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  "Text": "['Corrosion Science 170 (2020) 108645  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Long time ablation behaviors of designed ZrC-SiC-TiC ternary coatings for environments above 2000 °C  T  Xiaohui Pana,b, Yaran Niua,*, Xueting Xua, Xin Zhonga, Minhao Shia, Xuebin Zhenga,*, Chuanxian Dinga  a Key Laboratory of Inorganic Coating Materials CAS, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, 100049, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: ZrC-SiC-TiC ternary coatings Ablation resistance Oxide products Vacuum plasma spray  Thermal protection systems with long time anti-ablation property for applications in ultrahigh temperature environments have been a choke point for the development of hypersonic vehicles. In the present work, new ZrCSiC-TiC ternary composite coatings were designed and fabricated by vacuum plasma spray. The ablation behaviors of the ternary coatings were evaluated by a plasma flame with temperatures about 2200−2500 °C and compared with the ZrC-SiC and ZrC-TiC composite coatings. The characteristics of phases, element distributions and microstructure of the oxidized layers were observed in detail. The results showed that the anti-ablation resistant property of the ternary coatings was much better than those of the ZrC-SiC and ZrC-TiC composite coatings. The oxidation products, like high viscous SiO2 and high thermo stable TiO2, played a mutual promoted role in modifying the oxide layers, which contributing a lot to the improved ablation resistance. The ablation mechanisms of the ZrC-SiC-TiC ternary coatings were analyzed based on microstructure observations and thermodynamic analysis.  1.  Introduction  The severe operating conditions of hypersonic aircrafts such as ultra-high temperature, strong airflow scour and high thermal stresses have provoked urgent requirements for novel thermal protection systems (TPS) that have good oxidation resistance, thermal shock resistance and ablation resistance [1-4]. Ultra-high temperature ceramics (UHTCs) exhibit striking multi-characters, including extremely high melting points, high strength and modulus, relatively low density and so on, which are potential candidate materials for TPS [5-10]. ZrC has relatively low density (6.73 g/cm3), achigh melting point (3540 °C), ceptable resistance in ablation environment and low cost, which makes it attractive among all UHTCs. However, the application of mono ZrC is limited by its poor oxidation resistance. ZrC is susceptible to oxidation when temperature rises above 600 °C, forming porous ZrO2 scale that provide channels for oxygen diffusion [11-15]. Over the past years, SiC has been a major choice to improve the anti-ablation property of ZrC [16-21]. Jia et al. [16] fabricated ZrC20 vol.%SiC coating on C/C composites and investigated its ablation behavior usiing oxyacetylene flame. They found that after 120 s ablation, the linear ablation rate reduced by 96.4 % and the mass gain rate  increased by 383.3 % compared with those of the pure ZrC coating. The reports of ablation behaviors of ZrC-SiC systems were summarized in Table 1. It can be concluded that ZrC-SiC exhibited good ablation resistance in environments below 2200 °C for short time. The incorporation of SiC into ZrC could form a protective ZrO2-SiO2 oxide scale. SiO2 could serve as an ideal oxygen diffusion barrier because of its low for example, 3.2 × 10−15 gcm−1s−1 at 1000 °C oxygen permeability, [22]. However, the effect of SiC is limited by the environmental temperature for the following reasons: (i) the high evaporation and decomposition trend; (ii) the active oxidation of SiC inside. The evaporation or decomposition of SiO2 would leave a porous oxide scale, providing channels for oxygen diffusion. Besides, the active oxidation of SiC would generate large amount of gaseous products and cause damages. In the previous work in our laboratory, we designed and fabricated ZrC-TiC composite coatings with different TiC contents (20, 30, 40 vol. %) using vacuum plasma spray [23]. The ablation behaviors of this kind of coating were investigated under a plasma flame. The results showed that the ZrC-30 vol.%TiC coating exhibited better ablation resistance than that of the ZrC-30 vol.%SiC coating after 150 s ablation. The ZrC30 vol.%TiC coating kept integrated while the ablation center of ZrC ⁎ Corresponding authors. E-mail addresses: yrniu@mail.sic.ac.cn (Y. Niu), xbzheng@mail.sic.ac.cn (X. Zheng).  https://doi.org/10.1016/j.corsci.2020.108645 Received 6 September 2019; Received in revised form 23 March 2020; Accepted 25 March 2020  Available online 04 April 2020 0010-938X/ © 2020 Elsevier Ltd. All rights reserved.  \\x0c', 'X. Pan, et al.  Table 1  Summary of ablation behaviors of ZrC-SiC systems.  System  ZrC-20 vol.%SiC ZrC-20 vol.%SiC ZrC-SiC ZrC-SiC ZrC-30 vol.%SiC ZrC-SiC ZrC-SiC  System type  Ablation time  Ablation temperature  Coating Coating Coating Coating Ceramic Coating Coating  120 s 120 s 5−45 s 60 s 600 s 60 s 20 s  1738 °C 2105 °C 2173 °C -  2100 °C 2100 °C -  Corrosion Science 170 (2020) 108645  Results  Integrated Damaged Integrated Integrated Damaged Integrated Integrated  Refs  [16] [16] [17] [18] [19] [20] [21]  30 vol.%SiC peeled off. Based on the microstructure and thermal dynamic analysis, it was found that (Zr, Ti)O2 and TiO2 are more thermodynamically stable than SiO2, which contributed to the improved ablation resistance. However, we also found that the ZrC-30 vol.%TiC coating was almost totally oxidized after 150 s ablation. This phenomenon was attributed to the relatively high oxygen diffusion coefficients (1.12 × 10−13 m2s−1, of both TiO2 at 1800 °C) and ZrO2 (1.16 × 10−12 m2s−1, at 1800 °C) [24]. Things could become interesting if the advantages of SiO2 (low oxygen permeability) and TiO2 (high thermo stability) could be combined. However, the related works have not been reported yet. In the present work, ternary ZrC-SiC-TiC coatings were designed and prepared by vacuum plasma spray. The ablation behavior of two kinds of ternary composite coatings (ZrC-20 vol.%SiC-10 vol.%TiC, ZrC10 vol.%SiC-20 vol.%TiC) were evaluated by a plasma flame at atmosphere. ZrC-SiC (30 vol.%) and ZrC-TiC (30 vol.%) composite coatings were also prepared for comparison. It is expected that the SiC and TiC additions could play a positive joint function on the anti-ablation property of the ZrC coating.  2. Experimental  2.1. Coating preparation and characterization  Powder mixtures with different compositions, including ZrC-30 vol. %SiC (denoted ZS3), ZrC-20 vol.%SiC-10 vol.%TiC (denoted ZS2T1), ZrC-10 vol.%SiC-20 vol.%TiC (denoted ZS1T2) and ZrC-30 vol.%TiC (denoted ZT3), were prepared by mixing ZrC (Changsha Langfeng Metal Materials Co., Ltd., China), SiC (Qinhuangdao Yinuo Advanced Material Co., Ltd., China) and TiC (Qinhuangdao Yinuo Advanced Material Co., Ltd., China) by ball mixing in alcohol. Then, the powder mixtures were agglomerated by spray drying and sintered to obtain powders that were (Φ 40 mm × 5 mm) suitable for vacuum plasma spray. Graphite (Beijing Jinglong Carbon Materials Co., Ltd., China) with a SiC bonding layer was used as substrates. The preparation details of the SiC bonding layer could be seen elsewhere [25]. Vacuum plasma spray (A-2000, Sulzer Metco AG, Switzerland) was used to fabricate the coatings.  Hydrogen and argon were used to generate the plasma flame. The details of coating preparation could be seen in our previous work [26].  2.2. Ablation test  Argon and hydrogen were used to generate the plasma flame with a heat flux of 3.01 MW/m2 for evaluating the ablation resistance of the composite coatings. The heat flux was measured by a Gardon heat flux meter (Shanghai Tuxin electronic technology co., Ltd., China). All the ablation resistance tests were conducted at atmosphere for 300 s and 600 s, respectively. It should be stated that there was a certain time (about 30 s) for the plasma flame generator to cool down after a duration of 300 s during the 600 s ablation test. The coatings after 300 s and 600 s ablation were labeled as, taking the ZS3 coating for example, ZS3-300, ZS3-600, respectively. The temperatures of the ablation centers of the coating samples were recorded by a two-colour infrared pyrometer (Marathon MR1SC, Raytek, USA).  2.3. Composition and microstructure characterization  The phase compositions of the as-sprayed and ablated coatings were characterized by X-ray diffraction with Cu Ka (λ → 1.54056 Å) radiation (XRD, RAX-10, Rigaku, Japan). The step size and scanning rate are 0.005°and 10°/min, respectively. The microstructures and chemical element distributions of the as-sprayed and ablated coatings were characterized using a field emission scanning electron microscope (FEISEM, Magellen 400, USA) equipped with an energy dispersive spectrometer (EDS, PN-5502, INCA ENGERY, UK). The ablation centers of the coatings were chosen to evaluate the ablation resistance and take detailed microstructrue characterizations. The particle size of powders was evaluated by a laser particle analyzer (BT-9300 s, China). The porosity of the as-sprayed coatings was evaluated by three cross-sectional images with a magnification of 1000× using an image analysis software (Leica Qwin, Germany).  Fig. 1. Morphologies and particle size distribution of ZS3 starting powders.  2  \\x0c', 'X. Pan, et al.  Fig. 2. XRD patterns of as-sprayed ZS3, ZS2T1, ZS1T2 and ZT3 coatings.  3. Results  3.1. Phase compositions and microstructures of as-sprayed coatings  Fig. 1 shows the typical morphologies and particle size distributions of the ZS3 feedstock powders. It can be seen that the powder was (D50) was about 35−40 μm spherical and the measured mean size (Fig. 1b). Fig. 2 displays the XRD patterns of the as-sprayed coatings. The coatings were mainly composed of cubic ZrC. SiC could be observed in the ZS3 coating, but not in the ZS2T1 and ZS1T2 coatings, which could be ascribed to the relatively low content of SiC for those coatings. The positions of ZrC peaks were shifted about 0.5° higher, for the solid solution phenomenon of ZrC and TiC [27]. Fig. 3 shows the surface and cross-sectional morphologies of the as-sprayed coatings. Fully molten and partially molten particles could been seen in Fig. 3(a-d). No obviou cracks or voids could be observed between the SiC bonding layer and the top coating, which indicated good bonding for the coatings (Fig. 3(e-h)). The thickness of the ZS3, ZS2T1, ZS1T2 and ZT3 coatings were about 310, 360, 390 and 420 μm, respectively. The porosity of the composite coating was calculated using the crosssection SEM images, which were about 5-8 %.  3.2. Ablation behavior  Fig. 4 shows the macrographs of the coatings after 300 s and 600 s ablation tests. It can be seen that the morhologies of all coatings changed after the ablation tests and the ablation time had a significant influence. Only part proportion (about 40 %) of ZS oxide scale was retained after 300 s ablation (Fig. 4a), while the ZS2T1, ZS1T2 and ZT3  Corrosion Science 170 (2020) 108645  (Fig. 4b-d) coatings kept integrated (100 %) and seemed still bonded well to the substrate. When the ablation time was extended to 600 s, the unpeeling-off part of ZS coating is further decreased to 10 %, as shown in Fig. 4e (Yellow circle area). With the increasing of the TiC content, the ablated coating become more integrated. However, the ZT3 coating peeled off after 600 s ablation (Fig. 4h). The temperature/time curves of the coating samples during the ablation tests are illustrated in Fig. 5. The final temperature was decreased with the increasing of the TiC content, which were 2390 °C, 2363 °C, 2300 °C and 2266 °C, for the ZS, ZS2T1, ZS1T2 and ZT3 coating, respectively. The XRD patterns of the ablated coatings are shown in Fig. 6. For the coatings after 300 s ablation, it can been seen that dominant phase of all the coatings was monoclinic zirconia (m-ZrO2) (Fig. 6a). A trace of (Zr, Ti)O2 could be found in the ablated ZT3 coatings. (Zr, Ti)O2 was generated through reaction 1.  ZrO2(s) + TiO2(l) → (Zr, Ti)O2(l)  (1)  After 600 s ablation, phase transformation could be figured in Fig. 6b. That is, the intensity of m-ZrO2 dramatically declined and the peaks of (Zr, Ti)O2 became more distinct in all the TiC contained coatings compared with the 300 s ablated coatings. SiC could be found for the ZT3-600 coating (sticking to the substrate), indicating that the ZT3 coating was totally oxidized after 600 s ablation.  3.3. SEM observation  Representative surface and cross-sectional morphologies of all the coatings after 300 s ablation are shown in Fig. 7. The surface of the ZS3300 coating was composed of melted and solidified SiO2 and solid ZrO2 (Fig. 7a). SiO2 was also observed from the high magnification of the surface of the ZS2T1-300 and ZS1T2-300 coating (Fig. 7c and Fig. 7e). Besides, with the increasing of the TiC content, more melt area could be seen, especially for the ZT3-300 coating (Fig. 7g). The observation of the cross-sections of the ablated coatings indicated that a layered structure was formed after 300 s ablation for all the specimens. Combined with the structure and elemental mapping, different layers could be distinguished. Similarly, the bottom layers of all the ablated coatings were the un-oxidized layer for little O was detected. Therefore, the oxide scales were focused on. For the ZS3-300 coating, layer 1 was composed of glassy SiO2 and ZrO2 skeleton, presenting porous structure. And Layer 2 was supposed to be the SiC-depleted layer, which were confirmed by the Si mapping results (Fig. 7b). For the ZS2T1-300 coating, two layers were also observed for the oxide scale: layer 1 was similar to that of the ZS3-300 coating, which contained many pores. Layer 2 owned a relatively dense structure and almost no SiC-depleted area were detected. The oxide scale of the ZS1T2-300 coating was only composed of one layer, which seemed relatively dense (Fig. 7f). The microstructure of the ZT3-300 coating was similar to that of the ZS1T2300 coating. The cracks in Fig. 7(f-h) could be ascribed to the thermal  Fig. 3. Surface (a-d) and cross-sectional (e-h) morphologies of as-sprayed ZS3, ZS2T1, ZS1T2 and ZT3 coatings.  3  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 4. Macrographs of ZS3, ZS2T1, ZS1T2 and ZT3 coatings after ablation tests.  EDS results of the ZS3-600 coating. Almost no glassy SiO2 were observed which was confirmed by the EDS analysis (Fig. 8a). From the cross-sectional morphologies and elemental mappings, three distinct layers could be seen. Layer 1 and layer 2 belonged to the oxide scale and layer 3 was the un-oxidized scale. There was an crack in layer 2, which may be caused by the combination of the active oxidation of SiC and thermal stresses from the cooling process. The fractured morphologies of the ZS3-600 coating indicated that the structure of layer 1 was relatively porous and layer 2 was relative dense (Fig. 8(c-f)). Fig. 9 shows the surface, cross-section and related EDS results of the ZS2T1-600 coating. After 600 s ablation, more melted areas could be observed on the coating surface (Fig. 9a and b). The EDS results indicate that the melts were composed of Zr, Ti and O elements. From the cross-sectional morphology (Fig. 9c), it can be seen that the oxide scale could be divided into two layers. Layer 1 contained more pores compared with layer 2. The fractured morphologies of the ZS2T1-600 coating are given in Fig. 10. It can be seen that the area 1 of layer 1 was composed of particles associated with liquid phases (Fig. 10b). Layer 2 contained some threadlike pores (Fig. 10d) and layer 3 was the unoxidized layer (Fig. 10e). The surface, cross-section and related EDS results of the ZS1T2-600 coating are shown in Fig. 11. Almost no SiO2 could be seen on the surface (Fig. 11a). It is intersting to find that the oxide scale presented two distinct layers, which were different from that of the coating after  Fig. 5. Surface temperature curves of ZS3, ZS2T1, ZS1T2 and ZT3 coatings during ablation processes.  stresses during cooling process. The microstructure of the coatings after 600 s ablation were characterized in detail, trying to illustrate the effect of ablation time. Fig. 8 shows the surface, cross-sectional, fractured morphologies and related  Fig. 6. XRD patterns of ZS3, ZS2T1, ZS1T2 and ZT3 coatings after 300 s (a) and 600 s (b) ablation.  4  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 7. Surface, cross-sectional morphologies and related EDS results of ZS3-300 (a-b), ZS2T1-300 (c-d), ZS1T2-300 (e-f) and ZT3-300 (g-h) coatings.  300 s ablation. The difference of layer 1 and 2 might be resulted from the thermal gradient in the thick oxide scale. Moreover, there was a Ti rich area on the top surface, which exhibited a dense structure and might play as an effective oxygen diffusion barrier (Fig. 11c). The fractured morphologies of the ZS1T2-600 coating are illustrated in Fig. 12. It is worth mentioning that the crack in Fig.11c was caused by polishing because there was almost no crack can be observed from the fractured section (Fig. 12a). At area 1 in layer 1, particles associated with melts could be observed. Importantly, there was large amount of SiO2 retained in area 2, which was confirmed by the EDS result. Some  threadlike pores distributed in layer 2 (Fig. 12b) and layer 3 was the unoxidized coating (Fig. 12c). The thickness of the oxidized and un-oxidized layers of all the ablated coatings was measured and summarized in Table 2. It can be seen that the ZS1T2 coating exhibited the best oxidation resistance among all the coatings, having impact macrostructure and the largest thickness of the un-oxidized coating after 600 s ablation. From the above observations, the following phenomena are interesting and should be noted: (i) The ternary ZrC-SiC-TiC (ZS2T1 and ZS1T2) coating exhibited excellent ablation resistance compared with the ZS3 and ZT3 coatings;  Fig. 8. Surface (a), cross-sectional (b),  fractured morphologies (c-f) and related EDS results of ZS3-600 coating.  5  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 9. Surface (a-b), cross-sectional (c) and related EDS results of ZS2T1-600 coating.  (ii) the SiC-depleted layer was absent in the ablated ternary coatings.  4. Discussion  The ablation process is a combination of chemical corrosion and mechanical scour, leading to the loss of mass, the vaporization of liquid products and the escape of gaseous products. Several factors should be clarified before illustrating the failure mechanisms of the composite coatings: firstly, the difference in compositions, which would lead to different oxidation products; secondly, what is more important, the features and changes of the oxidation products during the ablation processes, like thermal stability, viscosity, oxygen permeability and so on. These factors were emphatically discussed as following.  4.1. Oxidation behaviors of ZrC, SiC and TiC  The chemical reactions of ZrC, SiC and TiC were influenced by both temperature and local oxygen pressure. For example, the oxygen partial pressure of layer 3 was significantly lower than that of the above layers. The following oxidation reactions may take place during the ablation processes [18,28]:  ZrC(s) + 2O2(g) → ZrO2(s) + CO2(g)  ZrC(s) + 3/2O2(g) → ZrO2(s) + CO(g)  SiC(s) + 2O2(g) → SiO2(l) + CO2(g)  SiC(s) + 3/2O2(g) → SiO2(l) + CO(g)  SiC(s) + O2(g) → SiO(g) + CO(g)  TiC(s) + 3/2O2(g) → TiO2(l) + CO(g)  TiC(s) + 2O2(g) → TiO2(l) + CO2(g)  TiC(s) + O2(g) → TiO(g) + CO(g)  (2)  (3)  (4)  (5)  (6)  (7)  (8)  (9)  4.1.1. Thermal dynamical prediction of oxidation products The volatility diagrams (the vapor pressure of predominant gaseous species, as a function of the equilibrium partial pressure of oxygen) of ZrC, SiC and TiC were constructed to predict the possible oxidation products [29,30]. The volatility diagrams of ZrC, SiC and TiC at 2300 °C were constructed by Factsage 7.2 software based on NIST-JANAF database [31], as shown in Fig. 13, aiming to predict the oxidation behavior of the ZS1T2 coating. The pO2 values for the generation of different condense phases of ZrC, SiC and TiC were calculated and summarized in Table 3. It can be seen that when 10−5 < pO2 < 100 Pa, there was a competition between active oxidation of SiC and oxidation of TiC. The XPS results of the inter layer of the ZS1T2-600 coating were supplied, as shown in Fig. 14. The XPS survey spectrum exhibited the presence of C, Zr, Si, Ti, and O elements (Fig. 14a). The high-resolution XPS spectrum of Ti 2p was well convoluted into four peaks (Fig. 14b): the binding energy at 458.1 eV and 464.1 eV was closed to that of Ti3+ 2p3/2 and Ti3+ 2p1/2 (457.9 eV, 464.0 eV), respectively [32,33]. While the other two peaks located at 459.1 eV and 464.9 eV were ascribed to Ti4+ 2p3/2 and Ti4+ 2p1/2 (458.9 eV, 464.5 eV), respectively. The presence of Ti3+ confirmed that there was some amount of Ti2O3 in the interlayer of the ZS1T2-600 coating. Combined with the absence of the SiC-depleted layer and the presence of Ti2O3, it could be induced that the competition of oxidation of TiC and active oxidation of SiC occurred during the ablation process, and the existence of Ti2O3 suppressed the formation of the SiC-depleted layer. A dense Ti-rich layer was observed in the ZS1T2-600 coating (Fig. 11c). Similar phenomenon had been observed by D.B. Lee [34], where a dense TiO2 layer was found for Ti3SiC2 after oxidation for 48 h in air at 1100 °C. The generation of the Ti-rich layer may be ascribed to the following two reasons: (i) TiO2 was less likely to evaporate or decompose, then retained to formed (Zr, Ti)O2 with ZrO2. While SiO2 was almost completely consumed after long time ablation. (ii) On the other hand, silicon ions in oxides could be relatively immobile, because of its  Fig. 10. Fractured morphologies of ZS2T1-600 coating.  6  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 11. Surface (a-b), cross-sectional (c) and related EDS results of ZS1T2-600 coating.  Fig. 12. Fractured morphologies and related EDS results of ZS1T2-600 coating.  Table 2  Thickness of oxide scale and un-oxidized coating at different ablation period.  Ablation time (s)  ZS3 (μm)  ZS2T1 (μm)  ZS1T2 (μm)  ZT3 (μm)  As-sprayed 300 600  Oxide  -  165 ± 10 295 ± 10  Un-oxidized  310 ± 10 210 ± 10 90 ± 15  Oxide  -  180 ± 5 390 ± 5  Un-oxidized  360 ± 10 255 ± 10 110 ± 10  Oxide  -  320 ± 5 340 ± 5  Un-oxidized  Oxide  Un-oxidized  390 ± 10 280 ± 10 180 ± 10  -  330 ± 10 Peeled  420 ± 10 290 ± 10 0  Fig. 13. Volatility diagrams of SiC (a), TiC (b) and ZrC (c) systems at 2300 °C.  higher bonding energy of Si+4-O (465 kJ mol−1) as compared with Ti+4-O (323 kJ mol−1) [34]. Therefore, it is possible for Ti4+ to gather in the outside layer and generate the Ti-rich layer. The consumption of SiO2 may be retarded because of the dense Ti-rich layer.  4.2. The features of SiO2 and TiO2  oxide scale and improving the oxygen diffusion resistance. Thermodynamic stability, viscosity and oxygen permeability are considered to be key factors for these phases, which would influence the ablation resistance of the coating and the microstructure of the oxide scale. The vapor pressure and viscosity of SiO2 and TiO2 were calculated and discussed as following:  ZrO2 was the skeleton of the oxide scale, and melted SiO2, TiO2 and (Zr,Ti)O2 played a role in associating the ZrO2 particles, densifying the  4.2.1. Thermodynamic stability During the ablation process,  SiO2  and TiO2 may  be  consumed  7  \\x0c', 'X. Pan, et al.  Table 3  Thermodynamic predictions of ZrC, SiC and TiC.  Corrosion Science 170 (2020) 108645  pO2 (Pa)  Predicted condense phase  Predicted gaseous phase  Predominant reaction  < 10−7  10−710−6 10−610−5 10−510−3  10−310° 100101 > 101  ZrC ZrC ZrC ZrC ZrO2 ZrO2 ZrO2 ZrO2  SiC SiC Si - - -  SiO2 SiO2  TiC TiC TiC TiC TiO Ti2O3 Ti3O5 TiO2  ZrC CO CO CO CO CO,CO2 CO2 CO2  SiC CO, Si, SiO CO, Si, SiO CO, SiO CO, SiO CO, SiO CO2, SiO2 CO2, SiO2  TiC CO, Ti, TiO CO, Ti, TiO CO, Ti, TiO CO, Ti, TiO CO, TiO CO2, TiO2 CO2, TiO2  - -  6 3, 6, 9 2, 3, 6 2, 4 2, 4, 8  through reaction (10−13).  SiO2(l) → SiO2(g)  SiO2(l) → SiO(g) +1/2O2(g)  TiO2(l) → TiO2(g)  TiO2(l) → TiO (g) +1/2O2(g)  (10)  (11)  (12)  (13)  The vapor pressure and decomposition pressure of SiO2 and TiO2 was calculated, as shown in Fig. 15. The vapor pressure and decomposition pressure of SiO2 were higher than that of TiO2, which meant SiO2 would be consumed easier during ablation processes. Besides, the calculated vaporization rate of SiO2 (207 mm/s) was nearly 900 times higher than that of TiO2 (0.23 mm/s) at 2225 °C [35], indicating a fast consuming of SiO2 in ultrahigh temperature environments.  4.2.2. Viscosity and oxygen permeability The viscosity of liquid phase has important influence on the mechanical scour resistance of the oxide scale. It is well known that viscosity as a function of temperature deviates from the Arrhenius relation [36]:  ln (η)  A  +  =  E RT  (14)  where η is the viscosity (dPa·s), A is the pre‐exponential factor (dPa·s), E is the activation energy (J mol−1), R is the gas constant and T is the absolute temperature. The viscosity of SiO2 and TiO2 were calculated by Factsage 7.2 using Viscosity module, as shown in Fig. 16. It can be seen that the viscosity of SiO2 was much higher than that of TiO2. Si-O tetrahedron could connect with each other and form Si-O network, while the network of Ti-O could hardly form at very high temperature [37]. Diffusivity of oxygen in liquid is inversely proportional to the viscosity of the liquid according to Stokes-Einstein relationship:  D=kT/6πηr  (15)  where D is diffusion constant [m2·s−1], k is Boltzmann’s constant [J·K−1], T is absolute temperature [K], η is viscosity [m2·s−1] and r is spherical particle radius [m]. It can be concluded that the oxygen diffusion rate in SiO2 was much lower than that in TiO2.  4.3. Ablation mechanism analysis of  the coatings  Distinguished ablation behaviors were observed for the ZS3, ZS2T1, ZS1T2 and ZT3 coatings in the present ablation conditions. The failure mechanisms of the four kinds of coatings could be concluded as follow:  4.3.1. For ZS3 coating SiO2 is characterized of low thermo stability, presenting high vapor and fast vaporization rate, which was tended to evaporated in a fast speed above 2000 °C. Therefore, the SiO2 in the ZS3 coating was consumed. The porous outer oxide layer would provide plenty of channels for oxygen diffusion. Besides, the active oxidation of SiC inside the coating was inevitable, and the significant generation of CO (g) and SiO (g) could lead to the creation of cracks/pores inside the coating (Fig. 8b), which caused the peeling of the oxide layers. Therefore, the ablation resistance of the ZS3 coating was not so good at environments above 2000 °C and the life-time is limited.  4.3.2. For ternary composite coating Due to the joint addition of TiC and SiC, the competition between the active oxidation of SiC and formation of Ti-contained phases helped a lot in avoiding the formation of SiC-depleted layer for both the ZS2T1 and ZS1T2 coatings. Thus, the ternary composite coatings maintained integrated. The variation of the contents of SiC and TiC additions had influence on the ablation behaviors. The consumption of SiO2 leaving porous oxide scale and make the ZS2T1 coating suffer severe oxidation  Fig. 14. XPS spectra of ZS1T2-600 coating: (a) survey spectra and (b) Ti 2p.  8  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 15. Evaporation pressures and decomposition pressures of SiO2 and TiO2 at different temperatures.  Data availability  All research data supporting this publication are directly available within this publication.  CRediT authorship contribution statement  Xiaohui Pan: Writing original draft, Writing review & editing. Yaran Niu: Project administration, Supervision, Writing review & editing. Xueting Xu: Resources, Data curation. Xin Zhong: Resources, Data curation. Minhao Shi: Methodology. Xuebin Zheng: Project administration, Writing review & editing, Supervision. Chuanxian Ding: Supervision.  Declaration of Competing Interest  Fig. 16. Viscosity of SiO2 and TiO2 at different temperatures.  during the long-time ablation process. While for the ZS1T2 coating, the generation of a Ti-rich layer on the coating surface reduced the consumption of SiO2, endowing this coating the best ablation resistance among all kinds of coatings.  4.3.3. For ZT3 coating The oxidation resistance of the ZT3 coating was the worst among all the composite coatings. Although the coating kept intact after 300 s ablation, due to the high oxygen permeability of ZrO2 and TiO2, the ZT3 coating was completely oxidized after 600 s oxidation.  5. Conclusions  ZrC-SiC-TiC ternary composite coatings were designed and fabricated, which exhibited much better long-time ablation resistance than the ZrC-SiC and ZrC-TiC coatings in environments above 2000 °C. The anti-ablation property of ZrC-10 vol.%SiC-20 vol.%TiC coating was the best under the present ablation condition. The competition between the oxidation of TiC and SiC partially avoided the generation of the SiCdepleted layer. The Ti-rich layer formed on the coating surface played an important role in resisting oxygen diffusion and SiO2 consumption, which contributed to the improved anti-ablation property of the ternary composite coating. This work convinced that the design of ternary composite coating would be a promising way to enhance the long-life ablation resistance of ZrC in the environment above 2000 °C.  9  they have no conflicts of  The authors declared that work. We declare that we do not have any commercial or associative interest that represents a conflict of interest in connection with the work submitted.  interest  to this  Acknowledgements  This work was supported by the National Natural Science Foundation (for Young Scholar) of China under Grant 51102267 and Youth Innovation Promotion Association CAS (2014223).  References  [3]  [2]  [1] R. Savino, L. Criscuolo, G.D. Di Martino, S. Mungiguerra, Aero-thermo-chemical characterization of ultra-high-temperature ceramics for aerospace applications, J. Eur. Ceram. Soc. 38 (2018) 2937-2953. S. Xuetao, L. Kezhi, L. Hejun, D. Hongying, C. Weifeng, L. 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Zhang, Graphene promoted triphasic N/Ti3+-TiO2 heterostructures: Iin-situ hydrothermal synthesis and enhanced photocatalytic performance, J. Alloys Compd. 785 (2019) 732-741. [34] D.B. Lee, S.W. Park, Oxidation of Ti3SiC2 between 900 and 1200 °C in Air, Oxid. Met. 67 (2006) 51-66. [35] A. Bronson, J. Chessa, An evaluation of vaporizing rates of SiO2 and TiO2 as protective coatings for ultrahigh temperature ceramic composites, J. Am. Ceram. Soc. 91 (2008) 1448-1452. [36] G. Zhang, K. Chou, K. Mills, A structurally based viscosity model Metall. Mater. Trans. B 45 (2014) 698-706. [37] W.H. Zachariasen, The atomic arrangement in glass, J. Am. Chem. Soc. 54 (1932) 3841-3851.  for oxide melts,  10  \\x0c']"
},{
  "_id": 3,
  "PDF": "A chemomechanical coupling model for oxidation and stress evolution in ZrB2–SiC.pdf",
  "Text": "['ARTICLE  A chemomechanical coupling model stress evolution in ZrB2-SiC  for oxidation and  Hailong Wang and Shengping Shena)  State Key Laboratory for Strength and Vibration of Mechanical Structures, School of Aerospace, Xi’an Jiaotong University, Xi’an 710049, China  (Received 19 September 2016; accepted 28 November 2016)  A chemomechanical coupling model is presented in the temperature range of 1200-1800 °C based on the microstructure during oxidation of ZrB2-SiC. The model includes the interaction of the oxidation rate and the mechanical stress. The stress is generated due to the constraint from the substrate to the lateral growth. The generated stress results in the shrink of the pores in the oxide. At the outer glassy layer surface, the boundary layer evaporation is adopted to describe the evaporation rate. Using the coupling model, the evolutions of the oxide layer thickness, weight gain, pore radius, and stress in both the oxide and substrate are provided, and the theoretical calculated results agree well with the reported experimental results. The results reveal large stress in the oxide layer during the oxidation process. By comparing the results of ZrB2 with different volume fractions of SiC, it is found that ZrB2 with higher volume fraction of SiC has more excellent oxidation resistance and smaller stress.  I.  INTRODUCTION  Transition metal diborides, especially ZrB2 and HfB2, have high melting temperature, high thermal conductivity and excellent oxidation resistance, which makes ultrahigh temperature ceramics (UHTCs) be the most promising candidates used in the re-entry and hypersonic vehicles as leading edge components.1-7 Moreover, within the family of UHTCs, ZrB2 has the lowest theoretical density (6.09 g/cm3), which is a desired property in aerospace field.8,9 However, when the temperature is higher than 1200 °C, the oxidation resistance of ZrB2 is poor, for the products of B2O3 volatilize fast, leaving the porous and nonprotective ZrO2 layer.10-13 Great efforts have been made to improve the ox ida t ion res is tance of ZrB2 a t over 1200 °C .3 ,14-17 The add i t ion of S iC can significantly improve the oxidation resistance at elevated temperature.15,18-21 The formed glassy borosilicate is more viscous and has lower evaporation pressure. Besides, it is a better barrier to oxygen diffusion than boria.22,23 So the glassy borosilicate can provide more efficient protection than boria. The oxidation mechanism of UHTCs is very complicated. Numerous investigators have performed a lot of experiments to observe and analyze the microstructure evolution and the oxidation products during oxidation of UHTCs and have a certain understanding on the oxidation mechanism.24-35 Wang et al.35 synthesized ZrC-SiC  powders and inves t iga ted the ox ida t ion behav ior . A t low tempera ture , ZrC qu ick ly ox id izes to form a nonprotec t ive ZrO2 layer , wh i le pro tec t ive S iO2 forms to improve the ox ida t ion res is tance a t ;1000 °C . Zhao e t a l .34 ana lyzed the m icros truc ture and phase compos i t ion for ZrC-30 vol% S iC ox ida t ion from low tempera ture to u l t rah igh tempera ture . L i e t a l .33 fabr ica ted three-d imens iona l need led C f /ZrC-SiC and s tud ied the mechan ica l proper t ies , m ic ros truc ture and h ightempe ra ture proper t ies . They conc luded tha t the exce l len t h ightempera ture proper t ies are due to the a l .31 forma t ion of ZrS iO4 on the surface . Zhang e t inves t iga ted the st ruc ture evo lu t ion and ox ida t ion behavior of ZrB2-S iC in a w ide tempera ture range . The SiC content significantly affects the structure evolution over 1800 °C while it has no apparent influence al.32 below 1600 °C. Brent Bargeron et studied the oxidation behavior of hafnium carbide in the temperature range 1400-2060 °C. Three distinct layers are found in the cross section, and the interlayer oxide is the best barrier for oxygen diffusion in the three layers. Lately, theore t ica l progress of UHTCs ox ida t ion has been ach ieved . Par thasara thy e t proposed severa l models to describe the oxidation progress. The model can calculate the weight gain and oxidation kinetics and the results are in reasonable agreement with the experimental data. However, the growth strain in the oxidation process is not considered. The growth strain accompanying the metal oxidation process has been known for a long time.38-41 Several theories have been proposed to explain the mechanism of the origin of the growth strain.42-45 The Clarke’s model is the most adopted. A lot of researches  a l .36 ,37  Contributing Editor: Yanchun Zhou a)Address all correspondence to this author. e-mail: sshen@mail.xjtu.edu.cn DOI: 10.1557/jmr.2016.489  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  Ó Materials Research Society 2017  1267  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4                                                \\x0c', 'are based on the Clarke’s model.46-51 Moreover, the stress induced by the oxidation process in return can affect the oxidation kinetics.52-57 Recently, based on the Parthasarathy et al.’s model,36 Zhou et al.58 proposed a thermo-chemo-mechanical model for the oxidation of ZrB2. Panicaud et al.59 reported the stress evolution in the case of oxidation of Zr alloy. The oxidation product of ZrB2-SiC is ZrO2, which is the same as that of Zr alloy. So there exists stress in the oxidation of ZrB2-SiC, too. The stress analysis in the oxidation process is very essential to predict the cracking and spalling of the oxide. However, there are few investigations focused on the chemomechanical coupling issue in the oxidation process of ZrB2-SiC. In th is work , a chemomechanica l mode l for the oxidation and stress evolution in ZrB2-SiC is proposed. The oxidation process can induce the stress and in return the stress can affect the molecular oxygen diffusion and oxidation reaction rate. The model can describe the evolutions of the stress in the oxide layer and the unoxidized substrate, oxide layer thickness, weight gain, pore radius, and pore volume fraction. The calculated results are consistent with the experimental observations.  II. Chemomechanical model  A. Model  framework  The morphology of the oxide scale, the oxidation products and phase present are derived from the literature oxidation schematic is shown in Fig. 1. Since the oxidation behavior of ZrB2-SiC is different in different temperature ranges, the oxidation temperature in the model is limited to 1200-1800 °C, which is the range of engineering interest. The substrate is ZrB2 containing dispersive SiC particles. The oxidation products are disconnected ZrO2 grains  studies.2,15,24,29,31,60 The  (in cross section) with a continuous porous region which is filled with glassy borosilicate. In the limited range of temperature in the work, the permeability of molecular oxygen through the ZrO2 grains is very small compared with that through the porous channels filled with glassy borosilicate, so the molecular oxygen permeating through the ZrO2 grains is neglected, as presented in Parthasarathy et al.’s work.36 The oxygen in the environment first dissolves in the glassy borosilicate then permeates through the outer glassy layer and the porous channels to react with the ZrB2 at the interface 2 (see Fig. 1). A portion of oxygen diffuses through the SiC-depleted region to react with SiC at the interface 1 (see Fig. 1). The amount of oxygen is assumed to be proportional to the volume fraction of SiC. The volume fraction of SiC in the substrate is fs. The porosity within ZrO2 region is fp. The glassy borosilicate volume fraction and the ZrO2 volume fraction fZrO2 in the region (2-3) of Fig. 1 are expressed as  fg  fg ¼ fs þ fp 1 \\x00 fs  ð  Þ  ;  ð1Þ  fZrO2 ¼ 1 \\x00 fp  \\x00  \\x01  1 \\x00 fs  ð  Þ ¼ 1 \\x00 fg  :  ð2Þ  B. One-dimensional model  At the temperature range of 1200-1800 °C, chemical reaction can be schematically expressed as  the  ZrB2 þ 5  2  O2 ¼ ZrO2 þ B2O3  ;  ð3Þ  SiC þ 3  2  O2 ¼ SiO2 þ CO  :  ð4Þ  The thickness evolutions of the oxide and outer glassy layer are expressed as (details refer to Appendix)  dhox  dt  ¼ VZrO2 fZrO2  k0  1 exp  gpVmrkk  RT  \\x12  \\x13 \\x10  C i2  O2  \\x11  5=2  ;  ð5Þ  dhg  dt  ¼  \"  k0  1 exp  gpVmrkk  RT  \\x12  \\x13 \\x10 \\x12  C i2  O2  \\x11  5=2 \\x00 JB2 O3\\x00vap  \\x0c\\x0c#  VB2 O3  þ  \"  5  3  fs  1 \\x00 fs  k0  1 exp  gpVmrkk \\x0c\\x0c \\x00 JSiO\\x00vap  RT  \\x13 \\x10  C i2  O2  \\x11  5=2  \\x00 JSiO2\\x00vap  \\x0c\\x0c#  VSiO2 \\x00 dhox  dt  fg  :  ð6Þ  Equations (5) and (6) are Eqs. (A24) and (A25) in the Appendix. The variables in Eqs. (5) and (6) are defined in Appendix. The recession rate of ZrB2 and SiC is related to the change rate of hox through molar volume and volume fraction.  FIG.  1. Oxidation schematic of ZrB2-SiC.  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1268  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c                                               \\x0c', 'dRZrB2  dt  ¼ VZrB2 VZrO2  fZrO2  ð1 \\x00 fs Þ  dhox  dt  ;  ð7Þ  dRSiC  dt  ¼ 5  3  VSiC fZrO2 VZrO2 1 \\x00 fs  ð  Þ  dhox  dt  ;  ð8Þ  where VZrB2 are the molar volume of ZrB2 and SiC, respectively. RZrB2 and RSiC are the recession thickness of ZrB2 and SiC, respectively. Then the SiC-depleted layer thickness hSiC is expressed as  and VSiC  easily  hSiC ¼ RSiC \\x00 RZrB2  :  ð9Þ  The weight gain due to the oxidation is expressed as  Wg ¼ hox \\x00 RZrB2 1 \\x00 fs  fZrO2 qZrO2 þ fgqg ÞqZrB2  \\x00  \\x01  þ hgqg \\x00 RSiC fsqSiC  ð  ;  ð10Þ  where qZrO2 , qSiC , qZrB2 , and qg are the density of ZrO2, SiC, ZrB2, and borosilicate, respectively. Accompanying the oxide layer growth, the stress is generated. The average stress in the oxide and substrate is expressed as follows (details refers to Appendix): Equations (11) and (12) are Eqs. (A48) and (A49) in  the Appendix. The definition of variables in Eqs. (11) and (12) refers to Appendix. For the porous oxide layer, the relationship between Yong’s modulus and porosity is61  Eox ¼  E0  ox  1 þ 4fg  1\\x00v2  ox  ð  Þ  p  ;  ð13Þ  where E0 no pore. The variation of pore radius r and borosilicate volume fraction fg is  ox  is the Yong’s modulus of oxide when there is  r ¼ r0 1 þ epore  \\x00 \\x00  \\x01 \\x01  ;  ð14Þ  fg ¼ f 0  g  1 þ epore  2  ;  ð15Þ  where r0 and f 0 g are the pore radius and glassy borosilicate volume fraction when there is no stress.  III. RESULTS AND DISCUSSIONS  The temperature is 1600 °C and the material properties used in the calculation are listed in Table I. Figure 2 shows the thickness calculated by the model and the experimental data. The total thickness (oxide 1 glass), the oxide thickness and the SiC-depleted region thickness are all plotted. The calculated total thickness and the oxide thickness are both in good agreement with the experimental data, which validates the proposed model. Since the SiC-depleted region is indistinct in experiment observation, the experimental data of SiC-depleted region  thickness is not available. So the calculated SiC-depleted region thickness is not compared with the experimental data. The solid line represents the total thickness of the oxide and glassy layer. The dotted line represents the thickness of the oxide. It can be easily seen from Fig. 2 that the outer glassy layer is almost the same thickness as the oxide thickness. The outer glassy layer is a barrier for oxygen diffusion which can provide excellent protection. So the good oxidation resistance of  _rox ¼  Box exp \\x00 Qox  RT  \\x12  \\x13  \\x00rox  ð  Þnox þ Bsub exp \\x00 Qsub  RT  \\x12  \\x13  \\x00roxhox=hsub  ð  Þnsub  \\x14  þ 1 \\x00 vox  E2  ox  _Eox \\x00 1 \\x00 vsub  Esub  _hoxhsub \\x00 hox  _hsub  h2  sub  \\x12  \\x13  rox \\x00 Dox  _hox  \\x15\\x1e \\x12  1 \\x00 vox  Eox  þ 1 \\x00 vsub  Esub  hox  hsub  \\x13  ;  ð11Þ  _rsub ¼  Box exp \\x00 Qox  RT  \\x12 \\x15\\x1e  \\x13  rsub hsub =hox  ð  Þnox þ Bsub exp \\x00 Qsub  RT  \\x12 \\x13  \\x13  rsub  ð  Þnsub þ  1 \\x00 vox  Eox  _hsub hox \\x00 hsub  _hox  h2  ox  \\x00 1 \\x00 vox  E2  ox  _Eox hsub  hox  \\x12  \\x13  rsub  \\x14  \\x00Dox  _hox  \\x00 1 \\x00 vox  Eox  hsub  hox \\x00 1 \\x00 vsub  Esub  \\x12  :  ð12Þ  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1269  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x1e \\x1e \\x1e                                               \\x0c', 'SiC-containing ZrB2 is due to relatively thick outer glassy layer. The weight gain calculated by the proposed model is compared with the experimental data in Fig. 3. The weight gain is controlled not only by the formation of oxidation products but also by the evaporation of the liquid products. The calculated weight gain is almost parabolic and in reasonable agreement with the experimental data. The stress evolution in the oxide layer is shown in Fig. 4. First, the compressive stress in the oxide layer increases with time, for the dominant factor is the growth strain. When the stress reaches the peak value, the stress begins to decrease with time, which is because the creep strain is dominant now. For different growth coefficient the trend of stress in the oxide layer is the same but value of the stress increases with increasing the  the  formation  of  the  Dox,  the  TABLE I. Material properties used in the calculation.  ZrO2  ZrB2-SiC  15% SiC  20% SiC  Young’s  modulus  (E/GPa)  115 (Ref. 62)  278 (Ref. 63)  283 (Ref. 63)  Poisson’s  ratio (m)  0.33 (Ref. 64)  0.16 (Ref. 65)  0.16 (Ref. 65)  Norton stress  exponent  (n)  ;1.3 (Ref. 66)  ;1 (Ref. 67)  ;1 (Ref. 67)  Material creep  coefficient [B/(MPa\\x00n/s)] Activation energy  1 \\x02 105 (Ref. 66)  0.0455 (Ref. 67)  0.2537 (Ref. 67)  [Q/(kJ/mol)]  517 (Ref. 66)  240 (Ref. 67)  258 (Ref. 67)  FIG. 2. Calculated oxide layer thickness compared with experimental data2,60 for ZrB2-20 vol% SiC oxidized at 1873 K [from Eqs. (5), (6), and (9)].  FIG. 3. Calculated weight gain compared with experimental data2,60 for ZrB2-20 vol% SiC oxidized at 1873 K [from Eq. (10)].  FIG. 4. Stress evolution in the oxide oxidized at 1873 K [from Eq. (11)].  layer  for ZrB2-20 vol% SiC  FIG. 5. Stress evolution in oxidized at 1873 K [from Eq.  the substrate (12)].  for  ZrB2-20  vol% SiC  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1270  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4                                                \\x0c', 'growth coefficient. Figure 5 shows the stress evolution in the substrate. The stress in the substrate is tensile. The stress increases first, then the stress almost remains the same. This is because the stress is dominated by the growth strain in the oxide layer in short oxidation time. As the oxidation time is prolonged, the creep strain becomes remarkable. The stress release rate is almost equal to the stress increasing rate. The stress almost remains unchanged. For different growth coefficient, the stress variation has the similar trend but increases with increasing the growth coefficient. Since the stress in the oxide layer can affect the pore radius and the borosilicate volume fraction, the pore radius and the borosilicate volume fraction varies with time too. Figures 6 and 7 show the pore radius evolution with time and the borosilicate volume fraction evolution with time, respectively. It can be seen that the pore radius  shrinks with time. The larger the growth coefficient is, the more the pore radius shrinks. The reason is that the compressive stress increases with the growth coefficient and the larger the compressive stress is, the more the pore radius shrinks. The borosilicate volume fraction has the same trend as the pore radius, for both of them are the characterizations of the effect of stress on the pore. Figures 8 and 9 show the effect of different volume fraction additions of SiC on the ZrB2-SiC oxidation behavior. It is obvious that the ZrB2 with 20% volume fraction addition of SiC has thinner oxide layer thickness and smaller weight gain than that with 15% volume fraction addition of SiC. The addition of SiC particles can significantly reinforce the oxidation resistance of ZrB2 at elevated temperature. This is because the products of glassy SiO2 have very small evaporation pressure. Besides, the diffusivity of oxygen in glassy SiO2 is  FIG. 8. Oxide layer thickness oxidized at 1873 K [from Eqs.  for ZrB2 with different SiC content (5), (6) and (9)].  FIG. 9. Weight gain for ZrB2 with different SiC contents oxidized at 1873 K [from Eq. (10)].  FIG. 6. Pore radius evolution for ZrB2-20 vol% SiC oxidized at 1873 K [from Eq. (14)].  FIG. 7. fg evolution for ZrB2-20 vol% SiC oxidized at 1873 K [from Eq. (15)].  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1271  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4                                                \\x0c', 'much smaller than that in glassy B2O3. The formed outer glassy borosilicate of ZrB2 with higher volume fraction addition of SiC contains more SiO2 than that with lower volume fraction addition of SiC. So the ZrB2 with higher volume fraction addition of SiC has more excellent oxidation resistance. The oxidation rate is slower, which results in thinner oxide layer thickness and smaller weight gain. Since the ZrB2 with 20% volume fraction addition of SiC has thinner oxide layer, it is obvious that the stress should also be smaller. It can be seen from Figs. 10 and 11 that both the stresses in the oxide layer and the substrate are smaller for 20% volume fraction addition of SiC.  IV. CONCLUSIONS  A chemomechanical model is proposed for the oxidation and stress evolution in ZrB2-SiC. The model  focuses on the coupling effects of chemical reaction and mechanics. The stress state in the oxidation process is very essential to predict the cracking and spalling of the oxide. Accompanying the oxidation process, the stress generates. The stress induced by the oxidation can affect the oxidation reaction rate and the oxygen diffusion. The volume fraction of SiC has great influence on the oxidation process. Due to that SiO2 has excellent barrier for oxygen diffusion and low evaporation pressure, the ZrB2 with higher volume fraction addition of SiC has thinner oxide layer thickness, lower weight gain and smaller stress. The calculated results are consistent with experimental data.  ACKNOWLEDGMENTS  This work is suppor ted by the Na t iona l Na tura l Sc ience Founda t ion of Ch ina (NSFC Grants No . 11632014 , 11372238 , 11302161 and 11302162) and the Chang J iang Scho lar program .  REFERENCES  1. K. Upadhya, J-M. Yang, and W.P. Hoffman: Materials ultrahigh temperature structural applications. Am. Ceram. Bull. 76(12), 51 (1997). 2. S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, and J.A. Salem: Evaluation of ultra-high temperature ceramics for aeropropulsion use. J. Eur. Ceram. Soc. 22(14), 2757 (2002). 3. A. Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby: Oxidation of ZrB2-SiC ceramics under atmospheric and reentry conditions. Refract. Appl. Trans. 1(2), 1 (2005). 4. 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Shen: Non-equilibrium thermodynamics and variational principles for fully coupled thermal-mechanical-chemical processes. Acta Mech. 224(12), 2895 (2013). J. Li: Chemical potential for diffusion in a Metall. 15(1), 21 (1981). 70. N. Swaminathan, J. Qu, and Y. Sun: An electrochemomechanical I. Theory. Philos. Mag. 87(11), theory of defects in ionic solids. 1705 (2007). 71. Y. Suo and S. Shen: General approach on chemistry and stress coupling effects during oxidation. J. Appl. Phys. 114(16), 164905 (2013). 72. O. Golan, A. Arbel, D. Eliezer, and D. Moreno: The applicability of Norton’s creep power law and its modified version to a singlecrystal superalloy type CMSX-2. Mater. Sci. Eng., A 216(1), 125 (1996). 73. L. Dormieux, D. Kondo, and (John Wiley & Sons, Chichester, U.K., 2006).  66.  67.  69.  stressed solid. Scr.  F-J. Ulm: Microporomechanics  APPENDIX: FORMULATION OF THE CHEMOMECHANICAL MODEL  1. Chemical Gibbs function variation  For the fully coupled chemomechanical process, the following chemical Gibbs function variation is obtained:68  dPgc ¼  Z  V  dgc dV þ dDgc \\x00 dQ\\x03 \\x00 dW \\x03 ¼ 0  ;  ðA1Þ  where gc is the chemical Gibbs function, and Dgc is the complement dissipative energy. Q\\x03 and W \\x03 are the complement heart absorbed from the external heart source and work of the external mechanical force. In the case of ZrB2- SiC oxidation at high temperature, the oxidation reaction occurs at the interface between the oxide and substrate. By the same derivational process in Ref. 68, one gets  dPgc ¼ 0 ¼  Z Z Z Z Z  a  rij nj \\x00 Ti  \\x03  \\x00 \\x00 Z Z Z  \\x01  dui da  \\x00  V  rij; j þ fi \\x00 q€ui  \\x01  dui dV \\x00  Z  C  rij nj  \\r\\rdui dC  \\x00  t  0  V  _cN þ cN vk ;k þ Ji;i  N  \\x00 \\x00 \\x08  \\x01  dlN dV ds  þ  t  0  a  Ji  N ni \\x00 JN  \\x03  \\x01  dlN da ds  \\x00  t  0  C  cN ðvk \\x00 Vk Þ  k  knk þ  Ji  N ni  \\x00vNr _wr gdlN dC ds  ;  ðA2Þ  where ri j , and q traction, displacement, body  Ti  \\x03 ,  u i,  f i,  are the force,  stress, and  surface density,  respectively. is the particle velocity. N , and lN are the molar fraction, diffusion flux, and chemical potential of species N, respectively. vNr and _wr are the stoichiometric coefficient and reaction rate of the rth reaction, respectively. is the diffusion flux of species N on the boundary. Vk is the velocity of the moving interface C. n i is the component of the normal \\x01k k denotes vector of the boundary and interface. the jump. Due to the arbitrariness of dui and dlN , one can get from Eq. (A2) the governing equations  vk  cN ,  Ji  JN  \\x03  rij; j þ fi ¼ q€ui  ðA3Þ  _cN ¼ \\x00cN vk ;k \\x00 Ji;i  N  ðA4Þ  and the boundary conditions  rijnj ¼ Ti N ni ¼ JN Ji  \\x03  on ar  \\x03  on adn  :  ðA5Þ  On the moving interface, one can get the jump conditions  rij nj  \\r\\r ¼ 0  ;  cN ðvk \\x00 Vk Þ  k  knk þ  Ji  N ni  \\r\\r \\x00 vNr _wr ¼ 0  :  ðA6Þ  2. One-dimensional model  At the temperature range of 1200-1800 °C, chemical reaction can be schematically expressed as  the  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1274  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\r\\r \\r\\r \\r\\r \\r\\r \\r\\r                                               \\x0c', 'ZrB2 þ 5  2  O2 ¼ ZrO2 þ B2O3  ;  ðA7Þ  SiC þ 3  2  O2 ¼ SiO2 þ CO  :  ðA8Þ  According to the law of mass reaction rate of Eq. (A7)  action,  the  chemical  is52,54  _w1 ¼ k1 C i2  O2  \\x10  \\x11  5=2  ;  ðA9Þ  where is the chem ica l reac t ion ra te cons tan t of Eq . (A7) . C i2 is the molar concen tra t ion of oxygen a t the in te rface 2 in F ig . 1 . The mo lar frac t ion and mo lar concen tra t ion are re la ted by where Vm is the average mo lar vo lume of per a tom . Based on the assump t ion tha t the amoun t of oxygen consumed by the S iC is propor t iona l to the vo lume frac t ion of S iC , the fo l low ing express ion is ob tained .  k1  O2  cO2 ¼ VmCO2 ,  the ox ide  _w2 ¼ 5  3  fs  1 \\x00 fs  k1 C i2  O2  \\x10  \\x11  5=2  ;  ðA10Þ  where is the chemical reaction rate of Eq. (A8). The generated ZrO2 causes the growth of the oxide layer. According to the conservation of zirconium atom, one can get  _w2  dhox  dt  ¼ VZrO2 fZrO2  _w1  ;  ðA11Þ  where hox and VZrO2 are the thickness of the oxide layer and molar volume of the ZrO2, respectively. The net increase of liquid products (SiO2 and B2O3) first fills up the pore in the oxide layer, then forms the outer glassy layer. The outer glassy layer thickness hg is expressed as  dhg  dt  ¼  _w1 \\x00 JB2 O3\\x00vap  \\x00  \\x01  VB2O3 \\x0c\\x0c \\x00 JSiO\\x00vap  þ  _w2 \\x00 JSiO2\\x00vap  \\x00  \\x01  VSiO2 \\x00 dhox  dt  fg  ;  ðA12Þ  where JB2O3\\x00vap , JSiO2\\x00vap , and JSiO\\x00vap are the evaporation flux of B2O3, SiO2, and SiO at the surface, respectively. VB2O3 and VSiO2 are the molar volume of B2O3 and SiO2, respectively. The recession rate of ZrB2 and SiC is related to the change rate of hox through molar volume and volume fraction.  dRZrB2  dt  ¼ VZrB2 VZrO2  fZrO2  1 \\x00 fs  ð  Þ  dhox  dt  ;  ðA13Þ  dRSiC  dt  ¼ 5  3  VSiC fZrO2 VZrO2 1 \\x00 fs  ð  Þ  dhox  dt  ;  ðA14Þ  where VZrB2 and VSiC are the molar volume of ZrB2 and SiC, respectively. and are the recession thickness of ZrB2 and SiC, respectively. Then the SiC-depleted layer thickness hSiC is expressed as  RZrB2  RSiC  easily  hSiC ¼ RSiC \\x00 RZrB2  :  ðA15Þ  The weight gain due to the oxidation is expressed as  Wg ¼ hox  fZrO2 qZrO2 þ fgqg  \\x00  \\x01  þ hgqg \\x00 RSiC fsqSiC ÞqZrB2  \\x00 RZrB2 1 \\x00 fs  ð  ;  ðA16Þ  where qZrO2 , qSiC , qZrB2 , and qg are the density of ZrO2, SiC, ZrB2, and borosilicate, respectively. Ignoring the convection item, Eq. (A4)  is written as  @CO2  @ t  ¼ \\x00 @ J O2  3  @ x3  ;  ðA17Þ  where J O2 is the flux of oxygen in x3 direction. The stress-dependent chemical potential of oxygen is69  3  lO2 ¼ l0  O2 þ RT ln VmCO2  ð  Þ \\x00 gO2 Vmrkk  ;  ðA18Þ  where R and T are the gas constant and absolute temperature, respectively. gO2 and rkk are the coefficient of compositional expansion and hydrostatic stress in the oxide layer. l0 is the chemical potential at the reference state. The oxygen flux is expressed as  O2  J O2  3  ¼ \\x00 DO2\\x00gCO2 RT  @lO2 @ x3 @CO2 @ x3  ¼ \\x00DO2\\x00g  \\x00 gO2 Vm RT  CO2  @rkk @ x3  \\x12  \\x13  ;  ðA19Þ  where DO2\\x00g borosilicate. Then, Eq.  is  the diffusivity of oxygen in the glassy  (A17) can be rewritten as  @CO2  @ t  ¼ \\x00 @ J O2  3  @ x3  ¼ DO2\\x00g  @ 2CO2  @ 2 x3  \\x00 gO2 Vm RT  @CO2 @ x3  @rkk @ x3  \\x12  \\x00 gO2 VmCO2  RT  @ 2rkk  @ 2 x3  \\x13  :  ðA20Þ  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1275  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c                                               \\x0c', 'Since the chemical potential is stress-dependent, the chemical reaction is stress-dependent too. The stressdependent chemical reaction rate constant is expressed as54,70  k1 ¼ k0  1 exp  gpVmrkk  RT  \\x12  \\x13  ;  ðA21Þ  where p represents the oxide. k0 is the chemical reaction rate constant when there is no stress. Stress-dependent chemical reaction rate is  1  _w1 ¼ k0  1 exp  gpVmrkk  RT  \\x12  \\x13 \\x10  C i2  O2  \\x11  5=2  ;  ðA22Þ  _w2 ¼ 5  3  fs  1 \\x00 fs  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2  :  ðA23Þ  Substituting Eqs. (A22) and (A23) into Eqs. (A11) and (A12), they are rewritten as  dhox  dt  ¼ VZrO2 fZrO2  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2  ;  ðA24Þ  dhg  dt  ¼  \"  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2 \\x00 JB2 O3\\x00vap  \\x0c\\x0c#  VB2O3  þ  \"  5  3  fs  1 \\x00 fs  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2  \\x00 JSiO2 \\x00vap  \\x0c\\x0c \\x00 JSiO\\x00vap  \\x0c\\x0c#  VSiO2 \\x00 dhox  dt  fg  :  ðA25Þ  The evaporation flux at the surface is express as follows  Jspecies\\x00vap  \\x0c\\x0c ¼ Dspecies  RT  105 pspecies\\x00vap dbdry  ;  dbdry ¼ 3  2  r ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi \\x12  lspecimen  vfluid  gfluid qfluid  \\x13  1=6  Dspecies  \\x00  \\x01  1=3  ;  ðA26Þ  where Dspecies  is the diffusivity in mixed gas. pspecies\\x00vap is the evaporation pressure. Species refers to SiO2, B2O3, and SiO. dbdry is the boundary layer thickness. is the length of the specimen. vfluid , gfluid , and qfluid are the velocity, viscosity, and density of the ambient fluid, respectively. According to Eq. (A6), the moving boundary condition the interface 2 is  lspecimen  of  J O2 ;i2  3  ¼ fgC i2  O2  V þ 5 2  _w1 þ 3  2  _w2  ¼ fgC i2  O2  VZrO2  fZrO2  k0  1 exp  gpVmrkk  RT  \\x12  \\x13 \\x13 \\x10  C i2  O2  \\x10  \\x11  5=2  þ  1 1 \\x00 fs  5  2  k0  1 exp  gpVmrkk  RT  \\x12  C i2  O2  \\x11  5=2  ;  ðA27Þ  where J O2 ;i2  3  is the oxygen flux at the interface 2. V is the the interface 2 and is expressed as  velocity of  V ¼ dhox dt  :  ðA28Þ  At  the interface 3,  the flux of  the oxygen is equal.  J O2 ;i3þ  3  ¼ J O2 ;i3\\x00 3  ðA29Þ  Since the quasi-steady state flux can be expressed as  is  assumed,  the oxygen  J O2 ;i3þ  3  ¼ D3a  Ca  O2 \\x00 C i3  O2  hg  ;  ðA30Þ  J O2 ;i3\\x00  3  ¼ D23  O2 \\x00 C i2 C i3  O2  hox  fg  ;  ðA31Þ  where D23 and D3a are the oxygen diffusivity in the region (2-3) and region (3-a), respectively. Since B2O3 evaporates much faster than SiO2 at the surface, the glassy layer contains more SiO2 than borosilicate filled in the porous channels. As a result, the diffusivity for oxygen in region (2-3) is different from that in region and Ca are the oxygen concentration at the interface 3 and a, respectively. Ca is expressed as  (3-a). C i3  O2  O2  O2  Ca  O2 ¼ pa  O2  S  ;  ðA32Þ  where pa and S are the oxygen partial pressure at interface and solubility of oxygen in the borosilicate. Substituting Eqs. (A30) and (A31) into Eq. (A29), we obtain  O2  the air glassy  C i3  O2 ¼ D3ahoxCa  O2 þ D23 fghgC i2 D23 fg hg þ D3ahox  O2  ;  ðA33Þ  and then we get  the oxygen flux at  the moving boundary  J O2 ;i2  3  ¼ D23 fg  O2 \\x00 C i2 C i3  O2  hox  :  ðA34Þ  Solving Eqs. (A27), (A33), and (A34) simultaneously, we obtain the oxygen concentration at the interface 2, . Accompanying the oxide layer growth, the stress is produced. Based on Clarke’s assumption,44,71 the lateral growth strain is proportional to the oxide layer thickness.  C i2  O2  eg  ox ¼ Doxhox  ;  ðA35Þ  where Dox is the lateral growth coefficient. The elastic strain is  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1276  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c                                               \\x0c', 'i ¼ 1 \\x00 vi ee  Ei  ri  ;  ðA36Þ  where can be replaced by ox for oxide and sub for substrate. mi , Ei , and ri are the Poisson ratio, Young’s modulus and plane stress, respectively. For the porous oxide layer, the relationship between Yong’s modulus and porosity is61  i  Eox ¼  E0  ox  1 þ 4fg  1\\x00v2  ox  ð  Þ  p  ;  ðA37Þ  where E 0 no pore. For the high-temperature oxidation, the creep effect should be taken into account. Here, Norton’s creep power law72 is adopted in the work.  ox  is the Yong’s modulus of oxide when there is  _ec  i ¼ signðri ÞBi exp \\x00 Qi  RT  \\x12  \\x13  ri  j  jni  ;  ðA38Þ  where B, Q, and are the material constant, activation energy for creep process and the stress exponent, respectively. Since the stress is compressive in the oxide layer and tensile in the substrate, signðrox Þ ¼ \\x001 and signðrsub Þ ¼ 1. The total strain in the oxide and substrate is expressed, respectively, as  n  eox ¼ eg  ox þ ee ox þ ec ox  ;  ðA39Þ  esub ¼ ee sub þ ec sub  :  ðA40Þ  For the cracking, bulking and spalling are not taken into account in the work, the strain in the oxide and substrate should be compatible.  eg  ox þ ee ox þ ec ox ¼ ee sub þ ec sub  ðA41Þ  The rate of Eq.  (A41)  is  ox þ _ee ox þ _ec ox ¼ _ee _eg  sub þ _ec sub  :  ðA42Þ  Substituting Eqs. Eq. (A42), we obtain  (A35),  (A36)  and  (A38)  into  Dox  _hox þ 1 \\x00 vox  Eox  _rox \\x00 1 \\x00 vox E 2  ox  _Eoxrox  \\x00Box exp \\x00 Qox  RT  \\x12  \\x13  \\x00rox  ð  Þnox ¼ 1 \\x00 vsub  Esub  _rsub  þBsub exp \\x00 Qsub  RT  \\x12  \\x13  rsub  ð  Þnsub  :  ðA43Þ  The force equilibrium equation is  Z  hox  rox dh þ  Z  hsub  rsub dh ¼ 0;  i:e:; roxhox  þ rsubhsub ¼ 0  ;  ðA44Þ  where hsub is the thickness of substrate. In the work, we deal with the average stress in the oxide and substrate. From Eq. (A44), rsub can be expressed as  rsub ¼ \\x00rox hox =hsub  :  ðA45Þ  The rate of Eq.  (A45)  is  _rsub ¼ \\x00 _rox hox =hsub \\x00 rox  _hox hsub \\x00 hox  _hsub  h2  sub  :  ðA46Þ  Substituting Eqs. obtain  (A45)  and (A46)  into Eq.  (A43), we  Dox  _hox þ 1 \\x00 vox  Eox  _rox \\x00 1 \\x00 vox E2  ox  _Eoxrox \\x00 Box exp \\x00 Qox  RT  \\x12  \\x13  \\x00rox  ð  Þnox  ¼ \\x00 1 \\x00 vsub  Esub  _rox hox =hsub þ rox  _hox hsub \\x00 hox  _hsub  h2  sub  \\x12 \\x12  \\x13  þBsub exp \\x00 Qsub  RT  \\x13  \\x00rox hox =hsub  ð  Þnsub  :  ðA47Þ  Rearranging Eq.  (A47), we get  _rox ¼  Box exp \\x00 Qox  RT  \\x12  \\x13  \\x00rox  ð  Þnox þ Bsub exp \\x00 Qsub  RT  \\x12  \\x13  \\x00rox hox =hsub  ð  Þnsub  \\x14  þ 1 \\x00 vox  E 2  ox  _Eox \\x00 1 \\x00 vsub  Esub  _hox hsub \\x00 hox  _hsub  h2 þ 1 \\x00 vsub  sub  \\x12  \\x13 \\x13  rox \\x00 Dox  _hox  \\x15\\x1e  1 \\x00 vox  Eox  Esub  hox  hsub  \\x12  :  ðA48Þ  In the same way, we get  _rsub ¼  Box exp \\x00 Qox  RT  \\x12  \\x13  rsub hsub =hox  ð  Þnox þ Bsub exp \\x00 Qsub  RT  \\x12  \\x13  rsub  ð  Þnsub  \\x14  þ 1 \\x00 vox  Eox  _hsub hox \\x00 hsub  _hox  h2  ox  \\x00 1 \\x00 vox  E2 hox \\x00 1 \\x00 vsub  ox  _E ox hsub  hox  \\x12  \\x13  rsub  \\x00Dox  _hox  \\x15\\x1e  \\x00 1 \\x00 vox  Eox  hsub  Esub  \\x12  \\x13  :  ðA49Þ  The equations can be calculated by the evolutionary algorithm by assuming an initial thickness of hox and hg . The variation of pore radius r and borosilicate volume fraction fg is  r ¼ r0  1 þ epore  \\x00 \\x00  \\x01 \\x01  ;  ðA50Þ  fg ¼ f 0  g  1 þ epore  2  ;  ðA51Þ  where r0 and f 0 are the pore radius and glassy borosilicate volume fraction when there is no stress.  g  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1277  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x1e \\x1e \\x1e                                               \\x0c', 'The relationship between the macroscopic strain and the microscopic strain in the pore is expressed as73  epore ¼ ðI \\x00 SÞ\\x001 : E  ;  ðA52Þ  where I is the forth-order identity tensor. S is the Eshelby tensor. E is the macroscopic strain. If the effect of lateral growth strain on the pore strain is neglected, the pore strain is  epore ¼ ðI \\x00 SÞ\\x001 :  ee  ox þ ec ox  \\x00  \\x01  :  ðA53Þ  To implement the model, the equations can be rewritten in incremental forms with respect to incremental time, Dt . Eqs. (A24) and (A25) are rewritten as  Dhox ¼ VZrO2 fZrO2  k 0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2  Dt  ;  ðA54Þ  Dhg ¼  \"  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2 \\x00 JB2 O3\\x00vap  \\x0c\\x0c#  VB2 O3 Dt  þ  \"  5  3  fs  1 \\x00 fs  k0  1 exp  gpVmrkk  RT  \\x12  \\x13  C i2  O2  \\x10  \\x11  5=2  \\x00 JSiO2\\x00vap  \\x0c\\x0c \\x00 JSiO\\x00vap  \\x0c\\x0c#  VSiO2 Dt \\x00 Dhox fg  :  ðA55Þ  The thickness at  t ðkþ1Þ can be easily obtained as  hðkþ1Þ  ox  ¼ hðkÞ  ox þ Dhox  ;  ðA56Þ  hðkþ1Þ  g  ¼ hðkÞ  g  þ Dhg  :  ðA57Þ  In the same way, Eqs. (A48) and (A49) are rewritten as  Drox ¼  Box exp \\x00 Qox  RT  \\x12  \\x13 \\x10 \\x11  \\x00rðkÞ  ox  \\x11 \\x12  nox Dt þ Bsub exp \\x00 Qsub  RT  \\x12  \\x13  \\x14 \\x10  \\x00rðkÞ  ox hðkÞ  ox  .  h  ðkÞ  sub  nsub Dt þ  1 \\x00 vox E2  ox  DEox  \\x00 1 \\x00 vsub  Esub  Dhoxh  ðkÞ  sub \\x00 h  ðkÞ  ox Dhsub  h  ðkÞ  sub  \\x10  \\x11  2  \\x13  rðkÞ  ox \\x00 DoxDhox  \\x15,  1 \\x00 vox  Eox  þ 1 \\x00 vsub  Esub  hðkÞ  ox  h  ðkÞ  sub  \\x12  \\x13  ;  ðA58Þ  Drsub ¼  Box exp \\x00 Qox  RT  \\x12 \\x12  \\x13 \\x10 \\x13 \\x10  rðkÞ  sub h  ðkÞ  sub  .  hðkÞ  ox  \\x11nox Dt  \\x14  þBsub exp \\x00 Qsub  RT  rðkÞ  sub  \\x11nsub Dt  þ  1 \\x00 vox  Eox  Dhsub h  ðkÞ  ox \\x00 h  ðkÞ  subDhox  h  ðkÞ  ox  \\x10  \\x11  2  \\x00 1 \\x00 vox  E2  ox  DEox h  ðkÞ  sub  hðkÞ  ox  0 B@  1  CArðkÞ  sub  \\x00DoxDhox  \\x15, \\x12  \\x00 1 \\x00 vox  Eox  h  ðkÞ  sub  hðkÞ  ox \\x00 1 \\x00 vsub Esub  \\x13  :  ðA59Þ  The stress at  t ðkþ1Þ can be easily obtained as  rðkþ1Þ  ox  ¼ rðkÞ  ox þ Drox  ;  ðA60Þ  r  ðkþ1Þ  sub  ¼ r  ðkÞ  sub þ Drsub  :  ðA61Þ  H. Wang et al.: A chemomechanical coupling model  for oxidation and stress evolution in ZrB2-SiC  J. Mater. Res., Vol. 32, No. 7, Apr 14, 2017  1278  D  o  w  n  l  o  a  d  e  d  f  r  o  m  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  .  F  l  o  r  i  d  a  I  n  t  e  r  n  a  t  i  n o  a  l  U  n  , i  n o  4 0  F  e  b  0 2 0 2  a  t  7 1  :  6 3  :  6 4  ,  s  b u  j  e  c  t  t  o  t  h  e  C  a  m  b  r  i  g d  e  C  o  r  e  t  e  r  m  s  o  f  u  s  e  ,  a  v  a  l i  a  b  l  e  a  t  h  t t  p  s  :  / /  w w w  .  c  a  m  b  r  i  g d  e  .  o  r  g  /  c  o  r  e  t /  e  r  m  s  .  h  t t  p  s  :  / /  o d  . i  o  r  g  /  0 1  .  7 5 5 1  /  j  m  r  .  6 1 0 2  .  9 8 4  \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x1e \\x1e \\x1e                                               \\x0c']"
},{
  "_id": 4,
  "PDF": "A thermoanalytical study on the oxidation of ZrC and HfC powders with formation of carbon.pdf",
  "Text": "['A thermoanalytical study on the oxidation of ZrC and HfC powders  with formation of carbon  Shiro Shimada *  Division of Materials and Engineering, Graduate School of Engineering, Hokkaido University, West-8,  North-13, Kitaku, Sapporo 060-8628 Japan  Received 30 August 2001;  received in revised form 13 March 2002; accepted 11 April 2002  Abstract  The oxidation of ZrC and HfC powders was  thermoanalytically investigated by simultaneous  thermogravimetry (TG),  differential thermal analysis (DTA) and mass spectrometry (MS) at various oxygen pressures ( PO2 the temperature range 20 - 1000 jC. TG results showed that the oxidation of ZrC and HfC begins at a fixed temperature of 380 and 400 jC, respectively, independent of PO2 oxide transformation, reached a maximum of 103 - 160%,  ) between 0.5 and 40 kPa in  and the type of sample. The degree of reaction, defined for the complete carbide -  then gradually returning to 100% at higher  temperatures. For  the  oxidation of ZrC, a sharp exothermic DTA peak appeared at PO2 10 kPa with the corresponding CO2 evolution but broad DTA and CO2 evolution peaks occurred at PO2 V 5 kPa. The oxidation of HfC gave two exothermic DTA peaks at all pressures, higher temperature peak agreeing with the CO2 evolution. Oxidation exceeding 100%, related with the formation of carbon,  z  the  is  discussed from the results of TG, DTA, MS and X-ray analysis. D 2002 Elsevier Science B.V. All  rights reserved.  Keywords: Oxidation; Formation of carbon; ZrC; HfC; TG; DTA; MS  1. Introduction  The oxidation of ZrC and HfC powders has been  investigated mainly from kinetic viewpoints by many  authors  [1 - 6].  It was pointed out  that  this  reaction  overshoots a degree of oxidation of 100% depending  on temperature  and oxygen pressure. Dufour  et  al.  [2]  and  Shimada  et  al.  [5,6]  suggested  that  this  overshoot ing  is  due  to  the  format ion  of  carbon  during oxidation. Barnier and Thevenot  [4]  reported  that  the  ox idation  of  zirconium oxycarbide  a lso  produces  a  high  carbon  content.  It  is  difficult  to  evidence the retention of carbon in the oxidation of  ZrC and HfC powders because of thinness in product  phase.  Thereafter, extensive studies have been performed  by Shimada et al.  [7 - 11,13] and Gozzi et al.  [12,14]  on the oxidation of ZrC, HfC and TiC using single  crystal in order to confirm the retention of carbon in a wide temperature range (500 - 1500 jC) at  relatively  low oxygen pressure (0.08 - 20 kPa) and to explain the  formation mechanism. The  former  authors  revealed  the formation of  two carbon-containing oxide scales,  dense  inside  and  porous  outside,  and  clarified  the  mechanism of carbon retention in the inner scale upon  oxidation of single crystals.  However, any attention has not been paid to the  evolution  of CO2  in  both  the  powder  and  crystal  0167-2738/02/$ see front matter D 2002 Elsevier Science B.V. All  rights reserved.  PII: S 0 1 6 7 2 7 3 8 ( 0 2 ) 0 0 1 8 0 7  * Fax: +81-11-706-6576.  E-mail address: shimashi@eng.hokudai.ac.jp (S. Shimada).  www.elsevier.com/locate/ssi  Solid State Ionics 149 (2002) 319 - 326  \\x0c', 'oxidation of  the carbides except  for  the oxidation of  NbC powder [13], which proceeds extremely rapidly jC by  above  485  a  grain-cracking  effect.  It  is  of  value  to  confirm the  formation  of  carbon without  oxidation by following the CO2 evolution during the  ox ida t ion  o f  Z rC and H fC .  The  p re sen t  s tudy  describes  the  oxidation  of ZrC and HfC powders  with formation of carbon from the measurements of  CO2  evolution,  combined with  simultaneous meas urements  of  both weight  and  thermal  changes.  It  aims  at  determining whether  the  evolution  of CO2  results  from the oxidation of ZrC or HfC or  that of  free  carbon  re ta ined  in  the  oxide  sca le w i thou t  oxidation.  2. Experimental  The starting ZrC and HfC powder  samples were  comme rc ia l ly  ava i lab le  (N ippon  Sh in -K inzoku ) ,  labelled N-ZrC and N-HfC,  respectively. The major  impurities  contained in N-ZrC and N-HfC powders and oxygen ( f 2.0  were  free  carbon ( < 0.5 wt.%)  wt.%). The grains of ZrC were of a flat shape, 1 - 10 Am in size and those of HfC were composed of aggregated particles of 2 - 5 Am in size with a main size of about 1 Am, as observed by SEM. The BET  surface area of the ZrC and HfC samples was 1.7 and \\x00 1, 1.0 m2 g  respectively.  The weight  and thermal  changes  and the  evolu tion  of CO2  during  oxidation were  simultaneously  monitored  by  thermogravimetry  (TG),  differential  thermal  analysis  (DTA)  and mass  spectrometry  (MS)  (TG - DTA 2000, MAC  Science; MS  (Q MS), VG Gas Analysis, Fisons  Instruments). About  10 mg of carbide was put in a Pt cell and heated at rate of 10 jC min \\x00 1 in a flowing O2 - Ar  a constant  gas mixture.  The  oxygen  partial  pressure  ( PO2  )  varied in a  range of 0.5 - 40 kPa by changing each  gas  flow rate. The total gas flow rate was \\x00 1. The phases  fixed to  be 100 ml min  formed were identi fied by X-ray analysis. For  a  comparison, ZrC and  HfC samples  of  different  origin  (Rare Metallic),  labelled R-ZrC and R-HfC, were  oxidized  under  the  same  conditions  as  above. Morphological  fea tures,  impurity contents  and surface  area of R-ZrC  and R-HfC powders are reported in previous papers  [5,6].  3. Results  3.1. Oxidation of ZrC powder  The oxidation of the N-ZrC samples was followed  at various PO2 TG - DTA and MS measurements, as shown in Figs. 1 and 2, respectively. The degree of oxidation, a (%),  between 1 and 40 kPa by simultaneous  in  the TG results, was  determined  by  dividing  the  observed weight  increase  by  the  theoretical  one,  which was calculated by assuming the complete con version of ZrC to ZrO2 according to the equation:  ZrC þ 2O2 ¼ ZrO2 þ CO2  ð1Þ  The TG results show that oxidation started at 380 jC  for any value of PO2 and was greatly accelerated for an a value around 70 - 80% at 590 - 600 jC at 40, 20 and  10 kPa (Fig. 1(A) - (C)). Oxidation passed on 100%  (see arrows), reached a maximum of 103 - 125%, and  th en w a s comp l e t ed by g r adu a l ly re tu rn ing a = 100% at 600, 700, 740 and 800 jC for PO2 = 40, 20, 10 and 5 kPa, respectively. The a value exceeding  to  100% was found to increase and slightly shift  to the  higher temperature with lowering PO2 . At PO2 oxidation slowly occurred, its degree of oxidation temperature (640 jC), went through 100% at higher then decreased down to 100% at 930 jC.  = 1 kPa,  An exothermic oxidation effect began around 400 jC at any value of PO2 . This effect formed a sharp DTA peak around 590 - 600 jC at 40, 20 and 10 kPa,  being consistent with the accelerated weight  increase  and became  less  important with decreasing PO2 contrast, a broad DTA peak was seen at 5 kPa, and at  .  In  1 kPa, two overlapping peaks were observed at 400 - 690 and 690 - 910 jC.  As shown by the MS curves (Fig. 2), the CO2 evolution began between 500 and 540 jC, depending . These values are 120 - 160 jC higher than the  on PO2  initiation temperature determined by TG and DTA, and  in these conditions,  the degree of oxidation is already  20 - 40%. The rapid CO2 evolution at 40 kPa coincides  with the rapid weight gain and the sharp DTA peak (Fig. 1(A) and (AV)). At 20 kPa, two overlapping CO2  evolution peaks were detected (Fig. 2(B)), the first peak corresponding to the sharp DTA peak (Fig. 2(BV)) and  the second peak to the gradual weight decrease above 600 jC. At 10 and 5 kPa (Fig. 2(C) and (D)),  the CO2  S. Shimada / Solid State Ionics 149 (2002) 319-326  320  \\x0c', 'S. Shimada / Solid State Ionics 149 (2002) 319-326  321  900 jC corresponding to the slow weight decrease above 700 jC (Fig. 1(E)) and to the second DTA peak above 690 jC (Fig. 1(EV)). The  temperature  corre sponding to the end of the CO2 evolution was consis tent with that  in the TG curve.  For the sake of comparison, the oxidation of R-ZrC  powder sample was performed at PO2  = 4, 10, 20 and  Fig. 1. Simultaneous TG - DTA curves in the oxidation of N-ZrC.  The solid lines correspond to TG and the dashed ones to DTA. PO2 (kPa) = 40 [(A) and (AV)], 20 [(B) and (BV)], 10 [(C) and (CV)], 5 (DV)], (EV)]. Heating jC min \\x00 1.  rate = 10  [(D)  and  1  [(E)  and  Sample mass = about 10 mg.  evolution formed a broad peak at 520 - 710 and 520 - 830 jC, respectively. At 1 kPa, the CO2 evolution (Fig. 2(E)) was rapid below 670 jC, at which temperature  the degree of oxidation has  reached 116%. Then,  it  Fig. 2. MS curves in the oxidation of N-ZrC. The marks (A) - (E)  slowly occurred over a wide temperature range of 670 -  and the oxidation conditions are the same as in Fig. 1.  \\x0c', '40 kPa and found to proceed in a very similar way to increase began at 380 jC that of N-ZrC. The weight and reached a maximum a value (104 - 132%) jC range, 570 - 720 gradually returning to 100% between 610 and 930 jC. The DTA curve showed a  in the  sharp exothermic peak for the oxygen pressure higher  than 10 kPa, but gave the lower and higher temperature peaks in the 400 - 700 and 700 - 940 jC range,  respectively,  at  4  kPa. For PO2 jC, at 530 - 590  z  10  kPa,  the CO2  evolution  began  i.e.  150 - 200  jC  higher  than the initiation temperature determined by  the weight and thermal changes, with a curve similar  to that of DTA. At 4 kPa, the CO2 evolution continued jC,  over  a wide  temperature  region  of  600 - 910  corresponding to the higher  temperature DTA peak.  Table  1  summarizes  the  initiation  temperatures the a  determined from TG, DTA and MS analysis,  values at the beginning of CO2 evolution, the maximum a values with the corresponding temperatures  and the completion temperatures for  the two types of  ZrC and HfC samples.  X-ray  ana lys is  of N-ZrC oxid ized  at  20  kPa  showed  that  the  crystalline phases of ZrO2 were jC but hardly formed for a = 30% at 500 slightly appeared for a = 100% at 580 jC. Oxidation going up  to 30% may be due to the formation of oxycarbide  ZrC1 \\x00 xOx or  amorphous ZrO2  [8],  the  latter being  crystallized to ZrO2 at 100% oxidation. Broad peaks  of monoclinic ZrO2 without ZrC appeared at the maximum degree of oxidation of 116% at 590 jC.  3.2. Oxidation of HfC powder  The oxidation of the N-HfC sample was carried out  in the PO2 TG - DTA and MS analysis, as shown in Figs. 3 and 4,  range of 0.5 - 40 kPa, with simultaneous  respectively. The  a  value was  also  determined  by  dividing the observed weight  increase by the theoret ical one calculated on the basis of  the equation:  HfC þ 2O2 ¼ HfO2 þ CO2  ð2Þ  From the weight tiated at 400 jC,  gain curves,  the  oxidation  ini independent of  the PO2 3). The weight increase was  value (see  dashed  lines  in Fig.  accelerated with  increasing  temperature,  passed  on  100% (see arrows) of f 160%,  and  reached  a maximum value  then returning to 100% at higher temper ature. At low PO2 , i.e. 1 and 0.5 kPa, the completion temperature increased up to 850 and above 1000 jC,  respectively.  Two separated exothermic DTA peaks appeared for  PO2  between 5 and 40 kPa, but these peaks overlapped  at 1 and 0.5 kPa (see dashed lines in Fig. 3). The lower temperature DTA peak began at 420 jC at any  PO2  , and the higher temperature one was found to shift  to higher  temperature with decreasing PO2 mum between the two exothermic peaks was located the temperature corresponding to the maximum a  . The mini at  value on the TG curves.  As displayed on the MS curves (Fig. 4), the CO2 evolution initiated at about 560 - 660 jC at all values is at a temperature 160 - 240 jC higher  of PO2  ,  that  than that determined by the TG measurements. In this  temperature range, oxidation has already proceeded to  as much as 50 - 100%, equivalent  to say that no CO2  evolution occurs before the degree of oxidation rea ches  these values.  In the 40 - 5 kPa range,  this CO2  evolution corresponds to the higher temperature DTA  Table 1  Initiation temperature, maximum oxidation and completion temper ature for ZrC and HfC samples  Sample  PO2  Initiation temperature (jC)  Maximum  Completion  TG  DTA MS/  oxidation  (%)  oxidation (%)/  temperature (jC)  temperature (jC)  N-ZrC  40  380  400  500/20  103/590  600  20  380  400  500/30  116/590  700  10  380  400  530/40  120/590  740  5  380  400  540/40  125/648 123/ f 740  800  1  380  400  540/40  930  R-ZrC  40  380  400  530/30  104/570  610  20  380  400  520/30  117/575  650  10  380  400  530/30  132/615  715  4  380  400  590/60  123/720  930  N-HfC  40  400  420  560/50  160/650  720  20  400  420  560/50  160/660  720  10  400  420  560/50  160/660  740  5  400  420  570/55  160/660  750  1  400  420  630/100  155/700  850  0.5  400  420  660/100  150/810  > 1000  R-HfC  20  400  420  560/15  102/700  800  10  400  420  560/16  113/710  850  4  400  420  580/26  120/780  880  1  400  420  600/30  130/790  940  0.5  400  420  680/60  122/880  >1000  S. Shimada / Solid State Ionics 149 (2002) 319-326  322  \\x0c', 'S. Shimada / Solid State Ionics 149 (2002) 319-326  323  Fig. 4. MS curves in the oxidation of N-HfC. The marks (A) - (F)  and the oxidation conditions are the same as in Fig. 3.  peak (Fig. 3(AV) - (DV)). At 1 and 0.5 kPa, tion of CO2 began at 630 and 660 jC,  the evolu respectively,  PO2  Fig. 3. Simultaneous TG - DTA curves in the oxidation of N-HfC. (kPa) = 40 [(A) and (AV)], 20 [(B) and (BV)], 10 [(C) and (CV)], 5 [(D) and (DV)], 1 [(E) and (EV)], 0.5 [(F) and (FV)]. Heating rate = 10 jC min \\x00 1. Sample mass = about 10 mg.  and was detected over a wide temperature interval of 260 - 300 jC.  The  oxidation of R-HfC powders was  also  per formed at PO2 = 0.5, 1, 4, 10 and 20 kPa and initiated at 400 jC (the same temperature as that of N-HfC) on  \\x0c', 'the TG curves. Oxidation exceeding 100% was also the maximum a  observed,  value  increasing  from  102% to 130% with PO2 to obtain the final 100% oxidation increased up to jC with PO2 kPa. The DTA curves  decreasing. The temperature  800,  850,  880,  940  and  above  1000  decreasing  from 20  to  0.5  showed only one broad, exothermic peak at PO2 and 10 kPa, but two separated and broad peaks at 4, 1  = 20  and 0.5 kPa. The CO2 evolution began at 560 - 680 jC, depending on PO2 has proceeded to 15 - 60%. The  , at which temperature oxidation  second DTA peak  coincided with that of CO2 evolution at PO2 0.5 kPa. These data are summarized in Table 1.  = 4, 1 and  X-ray analysis of N-HfC samples oxidized at 20  kPa showed that the crystalline phase of HfO2 was hardly formed even for a = 50% at 550 jC, but slightly appeared for a = 100% at 600 jC. Oxidation at 50%  may  be  re la ted  to  the  fo rma t ion  o f  oxyca rb ide  HfC1 \\x00 xOx or amorphous HfO2,  the latter being crys tallized to HfO2 at 100% oxidation. Broad peaks of  monoclinic HfO2 containing no HfC phase appeared at the maximum degree of oxidation of 160% at 650 jC.  4. Discussion  From the TG measurements,  the oxidation of ZrC jC,  and HfC begins  at  380  and  400  respectively,  independent of the value of PO2 (Table 1). The degree of oxidation passes on 100%  and the type of sample  and goes up to a maximum of 103 - 125% for N-ZrC  and  102 - 160% for N-HfC (Figs.  1  and  3).  If  the  carbon component of ZrC or HfC remains unoxidized  in the product,  the observed weight  increase would  equal 160% of  the theoretical one. Thus, exceeding  100% oxidation means that a considerable amount of  carbon  is  retained  in  the  product. This  effect was  found to increase with decreasing PO2 in previous papers [6 - 10], is associated with  , and as referred  the  formation of carbon according to Eqs.  (3) and (4):  ZrC þ O2 ¼ ZrO2 þ C  ð3Þ  HfC þ O2 ¼ HfO2 þ C  ð4Þ  Fig. 5 shows the equilibrium calculations for prod ucts (C, ZrO2, HfO2, CO, CO2) and reactants (ZrC, HfC) in the oxidation of ZrC or HfC at 600 jC, by using  the thermodynamic database HSC. This suggests that  the oxidation of ZrC or HfC at  low PO2 formation of equivalent amounts of ZrO2 or HfO2 and  occurs with  carbon, according to the reaction scheme of Eqs.  (3)  and (4). At higher PO2 CO2 (g). These equilibrium calculations  , carbon is oxidized to CO (g) or  support  the  assumption of a retention of carbon during the oxida tion of ZrC and HfC at  lower PO2 For N-ZrC, the initiation temperature of the CO2 evolution is 120 - 160 jC higher than that determined  values.  by TG and DTA (Figs. 1 and 2).  In these conditions,  the degree of oxidation is already about 20 - 40% with  formation  of ZrO2. At  10 - 40 kPa, the maximum about 600 jC corresponds  degree of 103 - 125% at  to the sharp DTA and CO2 evolution peaks. At 5 kPa,  the DTA peak becomes  less  sharp. As  reported pre viously [5],  this  abrupt  reaction may be due  to the  reaction of  fresh surfaces created by cracking of ZrC  grains  and  causing  the  rapid  oxidation  of  retained  carbon. At PO2 = 1 kPa, oxidation occurs slowly over a wide temperatures range of 360 jC giving two broad  overlapping DTA peaks, the latter one coinciding with  the CO2 evolution.  During the oxidation of N-HfC,  two DTA peaks  are also observed. They are separated at 5 - 40 kPa but  overlap each other at  lower pressure. The maximum  degree  a  of  160% corresponds  to  the  start  of  the  second DTA peak, which agrees with the initiation of  the CO2  evolution  (Figs.  3  and  4). XRD analysis  shows  that HfO2  forms  after  the  first DTA peak  without CO2  evolution. These  results  suggest  that  Fig. 5. Equilibrium calculations for moles of the products (C, ZrO2,  HfO2, CO, CO2) and the reactant (ZrC, HfC) on oxidation of ZrC or HfC at 600 jC.  S. Shimada / Solid State Ionics 149 (2002) 319-326  324  \\x0c', 'the first DTA peak is attributable to the formation of  HfO2 and carbon as represented by Eq. (4). A rapid weight decrease from a = 160% at higher temperatures  probably results from the oxidation of carbon remain ing in the HfO2  scale, giving the second DTA peak  accompanying the evolution of CO2.  It  is concluded  that the lower temperature DTA peak is due to the heat  generated by oxidation of Hf of  the HfC with for mation  of HfO2  and  carbon, whereas  the  higher  temperature DTA peak corresponds  to the evolution  of CO2 by oxidation of carbon remaining unreacted in  the HfO2 scale. At PO2 rates of oxidation of Hf and C are modified and the  lower  than 1 kPa,  the relative  two DTA peaks overlap each other.  The ratio of the area of the lower to the higher temperature DTA peak for N-ZrC (Fig. 1(AV) - (EV)) and N-HfC (Fig. 3(AV) - (EV)) is calculated to be about  2.2, which is near  the  ratio (2.5) of  the  formation  energy, at 800 K, of ZrO2 and HfO2 from Zr and Hf = \\x001091 and \\x001136 kJ mol \\x00 1, respectively) = \\x00 395 kJ mol \\x00 1). This to that of CO2 from C (DH800  (DH800  j  j  result supports the idea that  the low temperature DTA  peak is due to the oxidation of Zr or Hf and the high  temperature one to the burning of free carbon.  The initiation temperatures determined by TG and  MS  analysis  for  the  two  types  of ZrC and HfC  samples  are  summarized  in  Fig.  6. The  initiation  temperatures  of CO2 evolution are 100 - 300 and 5 symbols) determined from  jC  higher than that (  q  the TG measurements. At  these temperatures, oxida tion has already proceeded to 50 - 100%, with reten tion of carbon in the oxide scale. As PO2 the initiation temperature of CO2 evolution steeply  is lowered,  increases,  indicating that  lowering PO2 and increases the burning temper promotes reac tions  (3)  and (4)  ature of carbon retained in the oxide scale.  Raman  spec tra  of R-HfC sample a = 60% shows  oxid ized  a t  PO2  = 8  kPa  up  to  the  existence  of  amorphous  carbon  [6].  It  has  been reported that from about 550 jC  amorphous carbon is burning off  in air [11], which is near the initiation temperature (500 - 680 jC) of CO2 evolution (Table 1).  It can be  assumed that the initiation temperature of 380 - 400 jC observed by TG and DTA techniques  is  too low  for  oxidation  of  carbon  of ZrC or HfC,  but  high  enough for  the oxidation of Zr or Hf  into ZrO2 or  HfO2.  In our previous papers  [7 - 9],  it was demon strated  that  the  oxidation  of ZrC and HfC single  crystal  produces  two  oxide  scales:  an  inner  dense  layer  containing  23 - 24  at.% carbon  and  an  outer  porous one with 6 - 11 at.% carbon. Since the inner  scale acts as a barrier  for  the oxygen diffusion,  the  oxygen activity at  the interface between carbide and  inner scale becomes very low,  resulting in deposition  of carbon without oxidation.  In the oxidation of  the  ZrC and HfC powders,  such  an  oxygen  diffusion barrier  layer  of ZrO2  and HfO2  scales  probably  is  formed, contributing to a large decrease in the oxygen  activity  at  the  interface, which  deposits  elemental  carbon,  as  a  result  from the  data  in  Fig.  5. The  lowering of PO2 activity at the boundary, which is effective in deposi also leads to a decrease in the oxygen  tion of carbon from ZrC and HfC.  5. Conclusion  Simultaneous TG - DTA - MS analyses showed that  the  oxidation  of ZrC and HfC powders  at  various  Fig. 6. Dependency of the initiation temperature of CO2 on PO2 N-ZrC, .: R-ZrC, D: N-HfC, E: R-HfC. For a comparison, . initiation temperature of the weight increases by TG for Nand RZrC (5) and for Nand R-HfC (  o  :  the  q  ) are shown.  S. Shimada / Solid State Ionics 149 (2002) 319-326  325  \\x0c', '326  S. Shimada / Solid State Ionics 149 (2002) 319-326  oxygen pressures in the range 0.5 - 40 kPa begins at a fixed temperature of 380 - 400 jC, independent of PO2 and the type of sample. It reaches a maximum degree  of oxidation of 103 - 132% for ZrC and 102 - 160%  for HfC, depending on PO2 higher temperatures. Such  ,  then returning to 100% at  high  values  higher  than  References  [1] R.W. Bartlett, M.E. Wadsworth,  I.B. Cutler, Trans. Metall.  Soc. AIME 227 (1963) 467 - 472.  [2] L. Dufour, J. Simon, P. Barret, C. R. Acad. Sci., Ser. C 265  (1967) 171 - 174.  [3] A.K. Kuriakose,  J.L. Margrave,  J. Electrochem. Soc. 111  100% are attributable to the formation of  free carbon  (1964) 827 - 831.  in product phase.  Two exothermic DTA peaks  appeared at PO2 kPa for ZrC and in the range of 40 - 1 kPa for HfC. It  = 1  is explained that  the lower  temperature DTA peak is  [4] P. Barnier, F. Thevenot, Eur.  J. Solid State Inorg. Chem. 25  (1988) 495 - 508.  [5] S. Shimada, T.  Ishii,  J. Am. Ceram. Soc. 73 (1990) 2804 -  2808.  [6] S. Shimada, M.  Inagaki, K. Matsui,  J. Am. Ceram. Soc. 75  due to the oxidation of Zr or Hf component of ZrC or  (1992) 2671 - 2678.  HfC with formation of carbon and ZrO2 or HfO2 and  the higher  temperature DTA peak is associated with  the CO2 evolution by oxidation of  free carbon. This  assumption  is  also  supported  from thermodynamic  calculations.  Acknowledgements  The author wishes to express his hearty thanks to  Dr. M. Johnsson at University of Stockholm, Sweden,  for  thermodynamic calculations for oxidation of ZrC  and HfC.  [7] S. Shimada, M. Nishisako, M. Inagaki, J. Am. Ceram. Soc. 78  (1995) 41 - 48.  [8] S. Shimada, M.  Inagaki, M. Suzuki, J. Mater. Res. 11 (1996)  2594 - 2597.  [9] S. Shimada, K. Nakajima, M. Inagaki, J. Am. Ceram. Soc. 80  (1997) 1749 - 1756.  [10] S. Shimada, F. Yunazar, S. Otani,  J. Am. Ceram. Soc. 83  (2000) 721 - 728.  [11] S. Shimada, J. Ceram. Soc. Jpn. 109 (2001) S33 - S42(Special  Review).  [12] D. Gozzi, G. Cascino, S. Loreti, C. Minarini, S. Shimada, J.  Electrochem. Soc. 148 (2001) J15 - J24.  [13] S. Shimada, Oxid. Met. 42 (1994) 357.  [14] D. Gozzi, G. Guzzardi, A. Salleo, Solid State Ionics 83 (1996)  177 - 189.  \\x0c']"
},{
  "_id": 5,
  "PDF": "A thermoanalytical study on the oxidation of ZrC and HfC powders.pdf",
  "Text": "['A thermoanalytical study on the oxidation of ZrC and HfC powders  with formation of carbon  Shiro Shimada *  Division of Materials and Engineering, Graduate School of Engineering, Hokkaido University, West-8,  North-13, Kitaku, Sapporo 060-8628 Japan  Received 30 August 2001;  received in revised form 13 March 2002; accepted 11 April 2002  Abstract  The oxidation of ZrC and HfC powders was  thermoanalytically investigated by simultaneous  thermogravimetry (TG),  differential thermal analysis (DTA) and mass spectrometry (MS) at various oxygen pressures ( PO2 the temperature range 20 - 1000 jC. TG results showed that the oxidation of ZrC and HfC begins at a fixed temperature of 380 and 400 jC, respectively, independent of PO2 oxide transformation, reached a maximum of 103 - 160%,  ) between 0.5 and 40 kPa in  and the type of sample. The degree of reaction, defined for the complete carbide -  then gradually returning to 100% at higher  temperatures. For  the  oxidation of ZrC, a sharp exothermic DTA peak appeared at PO2 10 kPa with the corresponding CO2 evolution but broad DTA and CO2 evolution peaks occurred at PO2 V 5 kPa. The oxidation of HfC gave two exothermic DTA peaks at all pressures, higher temperature peak agreeing with the CO2 evolution. Oxidation exceeding 100%, related with the formation of carbon,  z  the  is  discussed from the results of TG, DTA, MS and X-ray analysis. D 2002 Elsevier Science B.V. All  rights reserved.  Keywords: Oxidation; Formation of carbon; ZrC; HfC; TG; DTA; MS  1. Introduction  The oxidation of ZrC and HfC powders has been  investigated mainly from kinetic viewpoints by many  authors  [1 - 6].  It was pointed out  that  this  reaction  overshoots a degree of oxidation of 100% depending  on temperature  and oxygen pressure. Dufour  et  al.  [2]  and  Shimada  et  al.  [5,6]  suggested  that  this  overshoot ing  is  due  to  the  format ion  of  carbon  during oxidation. Barnier and Thevenot  [4]  reported  that  the  ox idation  of  zirconium oxycarbide  a lso  produces  a  high  carbon  content.  It  is  difficult  to  evidence the retention of carbon in the oxidation of  ZrC and HfC powders because of thinness in product  phase.  Thereafter, extensive studies have been performed  by Shimada et al.  [7 - 11,13] and Gozzi et al.  [12,14]  on the oxidation of ZrC, HfC and TiC using single  crystal in order to confirm the retention of carbon in a wide temperature range (500 - 1500 jC) at  relatively  low oxygen pressure (0.08 - 20 kPa) and to explain the  formation mechanism. The  former  authors  revealed  the formation of  two carbon-containing oxide scales,  dense  inside  and  porous  outside,  and  clarified  the  mechanism of carbon retention in the inner scale upon  oxidation of single crystals.  However, any attention has not been paid to the  evolution  of CO2  in  both  the  powder  and  crystal  0167-2738/02/$ see front matter D 2002 Elsevier Science B.V. All  rights reserved.  PII: S 0 1 6 7 2 7 3 8 ( 0 2 ) 0 0 1 8 0 7  * Fax: +81-11-706-6576.  E-mail address: shimashi@eng.hokudai.ac.jp (S. Shimada).  www.elsevier.com/locate/ssi  Solid State Ionics 149 (2002) 319 - 326  \\x0c', 'oxidation of  the carbides except  for  the oxidation of  NbC powder [13], which proceeds extremely rapidly jC by  above  485  a  grain-cracking  effect.  It  is  of  value  to  confirm the  formation  of  carbon without  oxidation by following the CO2 evolution during the  ox ida t ion  o f  Z rC and H fC .  The  p re sen t  s tudy  describes  the  oxidation  of ZrC and HfC powders  with formation of carbon from the measurements of  CO2  evolution,  combined with  simultaneous meas urements  of  both weight  and  thermal  changes.  It  aims  at  determining whether  the  evolution  of CO2  results  from the oxidation of ZrC or HfC or  that of  free  carbon  re ta ined  in  the  oxide  sca le w i thou t  oxidation.  2. Experimental  The starting ZrC and HfC powder  samples were  comme rc ia l ly  ava i lab le  (N ippon  Sh in -K inzoku ) ,  labelled N-ZrC and N-HfC,  respectively. The major  impurities  contained in N-ZrC and N-HfC powders and oxygen ( f 2.0  were  free  carbon ( < 0.5 wt.%)  wt.%). The grains of ZrC were of a flat shape, 1 - 10 Am in size and those of HfC were composed of aggregated particles of 2 - 5 Am in size with a main size of about 1 Am, as observed by SEM. The BET  surface area of the ZrC and HfC samples was 1.7 and \\x00 1, 1.0 m2 g  respectively.  The weight  and thermal  changes  and the  evolu tion  of CO2  during  oxidation were  simultaneously  monitored  by  thermogravimetry  (TG),  differential  thermal  analysis  (DTA)  and mass  spectrometry  (MS)  (TG - DTA 2000, MAC  Science; MS  (Q MS), VG Gas Analysis, Fisons  Instruments). About  10 mg of carbide was put in a Pt cell and heated at rate of 10 jC min \\x00 1 in a flowing O2 - Ar  a constant  gas mixture.  The  oxygen  partial  pressure  ( PO2  )  varied in a  range of 0.5 - 40 kPa by changing each  gas  flow rate. The total gas flow rate was \\x00 1. The phases  fixed to  be 100 ml min  formed were identi fied by X-ray analysis. For  a  comparison, ZrC and  HfC samples  of  different  origin  (Rare Metallic),  labelled R-ZrC and R-HfC, were  oxidized  under  the  same  conditions  as  above. Morphological  fea tures,  impurity contents  and surface  area of R-ZrC  and R-HfC powders are reported in previous papers  [5,6].  3. Results  3.1. Oxidation of ZrC powder  The oxidation of the N-ZrC samples was followed  at various PO2 TG - DTA and MS measurements, as shown in Figs. 1 and 2, respectively. The degree of oxidation, a (%),  between 1 and 40 kPa by simultaneous  in  the TG results, was  determined  by  dividing  the  observed weight  increase  by  the  theoretical  one,  which was calculated by assuming the complete con version of ZrC to ZrO2 according to the equation:  ZrC þ 2O2 ¼ ZrO2 þ CO2  ð1Þ  The TG results show that oxidation started at 380 jC  for any value of PO2 and was greatly accelerated for an a value around 70 - 80% at 590 - 600 jC at 40, 20 and  10 kPa (Fig. 1(A) - (C)). Oxidation passed on 100%  (see arrows), reached a maximum of 103 - 125%, and  th en w a s comp l e t ed by g r adu a l ly re tu rn ing a = 100% at 600, 700, 740 and 800 jC for PO2 = 40, 20, 10 and 5 kPa, respectively. The a value exceeding  to  100% was found to increase and slightly shift  to the  higher temperature with lowering PO2 . At PO2 oxidation slowly occurred, its degree of oxidation temperature (640 jC), went through 100% at higher then decreased down to 100% at 930 jC.  = 1 kPa,  An exothermic oxidation effect began around 400 jC at any value of PO2 . This effect formed a sharp DTA peak around 590 - 600 jC at 40, 20 and 10 kPa,  being consistent with the accelerated weight  increase  and became  less  important with decreasing PO2 contrast, a broad DTA peak was seen at 5 kPa, and at  .  In  1 kPa, two overlapping peaks were observed at 400 - 690 and 690 - 910 jC.  As shown by the MS curves (Fig. 2), the CO2 evolution began between 500 and 540 jC, depending . These values are 120 - 160 jC higher than the  on PO2  initiation temperature determined by TG and DTA, and  in these conditions,  the degree of oxidation is already  20 - 40%. The rapid CO2 evolution at 40 kPa coincides  with the rapid weight gain and the sharp DTA peak (Fig. 1(A) and (AV)). At 20 kPa, two overlapping CO2  evolution peaks were detected (Fig. 2(B)), the first peak corresponding to the sharp DTA peak (Fig. 2(BV)) and  the second peak to the gradual weight decrease above 600 jC. At 10 and 5 kPa (Fig. 2(C) and (D)),  the CO2  S. Shimada / Solid State Ionics 149 (2002) 319-326  320  \\x0c', 'S. Shimada / Solid State Ionics 149 (2002) 319-326  321  900 jC corresponding to the slow weight decrease above 700 jC (Fig. 1(E)) and to the second DTA peak above 690 jC (Fig. 1(EV)). The  temperature  corre sponding to the end of the CO2 evolution was consis tent with that  in the TG curve.  For the sake of comparison, the oxidation of R-ZrC  powder sample was performed at PO2  = 4, 10, 20 and  Fig. 1. Simultaneous TG - DTA curves in the oxidation of N-ZrC.  The solid lines correspond to TG and the dashed ones to DTA. PO2 (kPa) = 40 [(A) and (AV)], 20 [(B) and (BV)], 10 [(C) and (CV)], 5 (DV)], (EV)]. Heating jC min \\x00 1.  rate = 10  [(D)  and  1  [(E)  and  Sample mass = about 10 mg.  evolution formed a broad peak at 520 - 710 and 520 - 830 jC, respectively. At 1 kPa, the CO2 evolution (Fig. 2(E)) was rapid below 670 jC, at which temperature  the degree of oxidation has  reached 116%. Then,  it  Fig. 2. MS curves in the oxidation of N-ZrC. The marks (A) - (E)  slowly occurred over a wide temperature range of 670 -  and the oxidation conditions are the same as in Fig. 1.  \\x0c', '40 kPa and found to proceed in a very similar way to increase began at 380 jC that of N-ZrC. The weight and reached a maximum a value (104 - 132%) jC range, 570 - 720 gradually returning to 100% between 610 and 930 jC. The DTA curve showed a  in the  sharp exothermic peak for the oxygen pressure higher  than 10 kPa, but gave the lower and higher temperature peaks in the 400 - 700 and 700 - 940 jC range,  respectively,  at  4  kPa. For PO2 jC, at 530 - 590  z  10  kPa,  the CO2  evolution  began  i.e.  150 - 200  jC  higher  than the initiation temperature determined by  the weight and thermal changes, with a curve similar  to that of DTA. At 4 kPa, the CO2 evolution continued jC,  over  a wide  temperature  region  of  600 - 910  corresponding to the higher  temperature DTA peak.  Table  1  summarizes  the  initiation  temperatures the a  determined from TG, DTA and MS analysis,  values at the beginning of CO2 evolution, the maximum a values with the corresponding temperatures  and the completion temperatures for  the two types of  ZrC and HfC samples.  X-ray  ana lys is  of N-ZrC oxid ized  at  20  kPa  showed  that  the  crystalline phases of ZrO2 were jC but hardly formed for a = 30% at 500 slightly appeared for a = 100% at 580 jC. Oxidation going up  to 30% may be due to the formation of oxycarbide  ZrC1 \\x00 xOx or  amorphous ZrO2  [8],  the  latter being  crystallized to ZrO2 at 100% oxidation. Broad peaks  of monoclinic ZrO2 without ZrC appeared at the maximum degree of oxidation of 116% at 590 jC.  3.2. Oxidation of HfC powder  The oxidation of the N-HfC sample was carried out  in the PO2 TG - DTA and MS analysis, as shown in Figs. 3 and 4,  range of 0.5 - 40 kPa, with simultaneous  respectively. The  a  value was  also  determined  by  dividing the observed weight  increase by the theoret ical one calculated on the basis of  the equation:  HfC þ 2O2 ¼ HfO2 þ CO2  ð2Þ  From the weight tiated at 400 jC,  gain curves,  the  oxidation  ini independent of  the PO2 3). The weight increase was  value (see  dashed  lines  in Fig.  accelerated with  increasing  temperature,  passed  on  100% (see arrows) of f 160%,  and  reached  a maximum value  then returning to 100% at higher temper ature. At low PO2 , i.e. 1 and 0.5 kPa, the completion temperature increased up to 850 and above 1000 jC,  respectively.  Two separated exothermic DTA peaks appeared for  PO2  between 5 and 40 kPa, but these peaks overlapped  at 1 and 0.5 kPa (see dashed lines in Fig. 3). The lower temperature DTA peak began at 420 jC at any  PO2  , and the higher temperature one was found to shift  to higher  temperature with decreasing PO2 mum between the two exothermic peaks was located the temperature corresponding to the maximum a  . The mini at  value on the TG curves.  As displayed on the MS curves (Fig. 4), the CO2 evolution initiated at about 560 - 660 jC at all values is at a temperature 160 - 240 jC higher  of PO2  ,  that  than that determined by the TG measurements. In this  temperature range, oxidation has already proceeded to  as much as 50 - 100%, equivalent  to say that no CO2  evolution occurs before the degree of oxidation rea ches  these values.  In the 40 - 5 kPa range,  this CO2  evolution corresponds to the higher temperature DTA  Table 1  Initiation temperature, maximum oxidation and completion temper ature for ZrC and HfC samples  Sample  PO2  Initiation temperature (jC)  Maximum  Completion  TG  DTA MS/  oxidation  (%)  oxidation (%)/  temperature (jC)  temperature (jC)  N-ZrC  40  380  400  500/20  103/590  600  20  380  400  500/30  116/590  700  10  380  400  530/40  120/590  740  5  380  400  540/40  125/648 123/ f 740  800  1  380  400  540/40  930  R-ZrC  40  380  400  530/30  104/570  610  20  380  400  520/30  117/575  650  10  380  400  530/30  132/615  715  4  380  400  590/60  123/720  930  N-HfC  40  400  420  560/50  160/650  720  20  400  420  560/50  160/660  720  10  400  420  560/50  160/660  740  5  400  420  570/55  160/660  750  1  400  420  630/100  155/700  850  0.5  400  420  660/100  150/810  > 1000  R-HfC  20  400  420  560/15  102/700  800  10  400  420  560/16  113/710  850  4  400  420  580/26  120/780  880  1  400  420  600/30  130/790  940  0.5  400  420  680/60  122/880  >1000  S. Shimada / Solid State Ionics 149 (2002) 319-326  322  \\x0c', 'S. Shimada / Solid State Ionics 149 (2002) 319-326  323  Fig. 4. MS curves in the oxidation of N-HfC. The marks (A) - (F)  and the oxidation conditions are the same as in Fig. 3.  peak (Fig. 3(AV) - (DV)). At 1 and 0.5 kPa, tion of CO2 began at 630 and 660 jC,  the evolu respectively,  PO2  Fig. 3. Simultaneous TG - DTA curves in the oxidation of N-HfC. (kPa) = 40 [(A) and (AV)], 20 [(B) and (BV)], 10 [(C) and (CV)], 5 [(D) and (DV)], 1 [(E) and (EV)], 0.5 [(F) and (FV)]. Heating rate = 10 jC min \\x00 1. Sample mass = about 10 mg.  and was detected over a wide temperature interval of 260 - 300 jC.  The  oxidation of R-HfC powders was  also  per formed at PO2 = 0.5, 1, 4, 10 and 20 kPa and initiated at 400 jC (the same temperature as that of N-HfC) on  \\x0c', 'the TG curves. Oxidation exceeding 100% was also the maximum a  observed,  value  increasing  from  102% to 130% with PO2 to obtain the final 100% oxidation increased up to jC with PO2 kPa. The DTA curves  decreasing. The temperature  800,  850,  880,  940  and  above  1000  decreasing  from 20  to  0.5  showed only one broad, exothermic peak at PO2 and 10 kPa, but two separated and broad peaks at 4, 1  = 20  and 0.5 kPa. The CO2 evolution began at 560 - 680 jC, depending on PO2 has proceeded to 15 - 60%. The  , at which temperature oxidation  second DTA peak  coincided with that of CO2 evolution at PO2 0.5 kPa. These data are summarized in Table 1.  = 4, 1 and  X-ray analysis of N-HfC samples oxidized at 20  kPa showed that the crystalline phase of HfO2 was hardly formed even for a = 50% at 550 jC, but slightly appeared for a = 100% at 600 jC. Oxidation at 50%  may  be  re la ted  to  the  fo rma t ion  o f  oxyca rb ide  HfC1 \\x00 xOx or amorphous HfO2,  the latter being crys tallized to HfO2 at 100% oxidation. Broad peaks of  monoclinic HfO2 containing no HfC phase appeared at the maximum degree of oxidation of 160% at 650 jC.  4. Discussion  From the TG measurements,  the oxidation of ZrC jC,  and HfC begins  at  380  and  400  respectively,  independent of the value of PO2 (Table 1). The degree of oxidation passes on 100%  and the type of sample  and goes up to a maximum of 103 - 125% for N-ZrC  and  102 - 160% for N-HfC (Figs.  1  and  3).  If  the  carbon component of ZrC or HfC remains unoxidized  in the product,  the observed weight  increase would  equal 160% of  the theoretical one. Thus, exceeding  100% oxidation means that a considerable amount of  carbon  is  retained  in  the  product. This  effect was  found to increase with decreasing PO2 in previous papers [6 - 10], is associated with  , and as referred  the  formation of carbon according to Eqs.  (3) and (4):  ZrC þ O2 ¼ ZrO2 þ C  ð3Þ  HfC þ O2 ¼ HfO2 þ C  ð4Þ  Fig. 5 shows the equilibrium calculations for prod ucts (C, ZrO2, HfO2, CO, CO2) and reactants (ZrC, HfC) in the oxidation of ZrC or HfC at 600 jC, by using  the thermodynamic database HSC. This suggests that  the oxidation of ZrC or HfC at  low PO2 formation of equivalent amounts of ZrO2 or HfO2 and  occurs with  carbon, according to the reaction scheme of Eqs.  (3)  and (4). At higher PO2 CO2 (g). These equilibrium calculations  , carbon is oxidized to CO (g) or  support  the  assumption of a retention of carbon during the oxida tion of ZrC and HfC at  lower PO2 For N-ZrC, the initiation temperature of the CO2 evolution is 120 - 160 jC higher than that determined  values.  by TG and DTA (Figs. 1 and 2).  In these conditions,  the degree of oxidation is already about 20 - 40% with  formation  of ZrO2. At  10 - 40 kPa, the maximum about 600 jC corresponds  degree of 103 - 125% at  to the sharp DTA and CO2 evolution peaks. At 5 kPa,  the DTA peak becomes  less  sharp. As  reported pre viously [5],  this  abrupt  reaction may be due  to the  reaction of  fresh surfaces created by cracking of ZrC  grains  and  causing  the  rapid  oxidation  of  retained  carbon. At PO2 = 1 kPa, oxidation occurs slowly over a wide temperatures range of 360 jC giving two broad  overlapping DTA peaks, the latter one coinciding with  the CO2 evolution.  During the oxidation of N-HfC,  two DTA peaks  are also observed. They are separated at 5 - 40 kPa but  overlap each other at  lower pressure. The maximum  degree  a  of  160% corresponds  to  the  start  of  the  second DTA peak, which agrees with the initiation of  the CO2  evolution  (Figs.  3  and  4). XRD analysis  shows  that HfO2  forms  after  the  first DTA peak  without CO2  evolution. These  results  suggest  that  Fig. 5. Equilibrium calculations for moles of the products (C, ZrO2,  HfO2, CO, CO2) and the reactant (ZrC, HfC) on oxidation of ZrC or HfC at 600 jC.  S. Shimada / Solid State Ionics 149 (2002) 319-326  324  \\x0c', 'the first DTA peak is attributable to the formation of  HfO2 and carbon as represented by Eq. (4). A rapid weight decrease from a = 160% at higher temperatures  probably results from the oxidation of carbon remain ing in the HfO2  scale, giving the second DTA peak  accompanying the evolution of CO2.  It  is concluded  that the lower temperature DTA peak is due to the heat  generated by oxidation of Hf of  the HfC with for mation  of HfO2  and  carbon, whereas  the  higher  temperature DTA peak corresponds  to the evolution  of CO2 by oxidation of carbon remaining unreacted in  the HfO2 scale. At PO2 rates of oxidation of Hf and C are modified and the  lower  than 1 kPa,  the relative  two DTA peaks overlap each other.  The ratio of the area of the lower to the higher temperature DTA peak for N-ZrC (Fig. 1(AV) - (EV)) and N-HfC (Fig. 3(AV) - (EV)) is calculated to be about  2.2, which is near  the  ratio (2.5) of  the  formation  energy, at 800 K, of ZrO2 and HfO2 from Zr and Hf = \\x001091 and \\x001136 kJ mol \\x00 1, respectively) = \\x00 395 kJ mol \\x00 1). This to that of CO2 from C (DH800  (DH800  j  j  result supports the idea that  the low temperature DTA  peak is due to the oxidation of Zr or Hf and the high  temperature one to the burning of free carbon.  The initiation temperatures determined by TG and  MS  analysis  for  the  two  types  of ZrC and HfC  samples  are  summarized  in  Fig.  6. The  initiation  temperatures  of CO2 evolution are 100 - 300 and 5 symbols) determined from  jC  higher than that (  q  the TG measurements. At  these temperatures, oxida tion has already proceeded to 50 - 100%, with reten tion of carbon in the oxide scale. As PO2 the initiation temperature of CO2 evolution steeply  is lowered,  increases,  indicating that  lowering PO2 and increases the burning temper promotes reac tions  (3)  and (4)  ature of carbon retained in the oxide scale.  Raman  spec tra  of R-HfC sample a = 60% shows  oxid ized  a t  PO2  = 8  kPa  up  to  the  existence  of  amorphous  carbon  [6].  It  has  been reported that from about 550 jC  amorphous carbon is burning off  in air [11], which is near the initiation temperature (500 - 680 jC) of CO2 evolution (Table 1).  It can be  assumed that the initiation temperature of 380 - 400 jC observed by TG and DTA techniques  is  too low  for  oxidation  of  carbon  of ZrC or HfC,  but  high  enough for  the oxidation of Zr or Hf  into ZrO2 or  HfO2.  In our previous papers  [7 - 9],  it was demon strated  that  the  oxidation  of ZrC and HfC single  crystal  produces  two  oxide  scales:  an  inner  dense  layer  containing  23 - 24  at.% carbon  and  an  outer  porous one with 6 - 11 at.% carbon. Since the inner  scale acts as a barrier  for  the oxygen diffusion,  the  oxygen activity at  the interface between carbide and  inner scale becomes very low,  resulting in deposition  of carbon without oxidation.  In the oxidation of  the  ZrC and HfC powders,  such  an  oxygen  diffusion barrier  layer  of ZrO2  and HfO2  scales  probably  is  formed, contributing to a large decrease in the oxygen  activity  at  the  interface, which  deposits  elemental  carbon,  as  a  result  from the  data  in  Fig.  5. The  lowering of PO2 activity at the boundary, which is effective in deposi also leads to a decrease in the oxygen  tion of carbon from ZrC and HfC.  5. Conclusion  Simultaneous TG - DTA - MS analyses showed that  the  oxidation  of ZrC and HfC powders  at  various  Fig. 6. Dependency of the initiation temperature of CO2 on PO2 N-ZrC, .: R-ZrC, D: N-HfC, E: R-HfC. For a comparison, . initiation temperature of the weight increases by TG for Nand RZrC (5) and for Nand R-HfC (  o  :  the  q  ) are shown.  S. Shimada / Solid State Ionics 149 (2002) 319-326  325  \\x0c', '326  S. Shimada / Solid State Ionics 149 (2002) 319-326  oxygen pressures in the range 0.5 - 40 kPa begins at a fixed temperature of 380 - 400 jC, independent of PO2 and the type of sample. It reaches a maximum degree  of oxidation of 103 - 132% for ZrC and 102 - 160%  for HfC, depending on PO2 higher temperatures. Such  ,  then returning to 100% at  high  values  higher  than  References  [1] R.W. Bartlett, M.E. Wadsworth,  I.B. Cutler, Trans. Metall.  Soc. AIME 227 (1963) 467 - 472.  [2] L. Dufour, J. Simon, P. Barret, C. R. Acad. Sci., Ser. C 265  (1967) 171 - 174.  [3] A.K. Kuriakose,  J.L. Margrave,  J. Electrochem. Soc. 111  100% are attributable to the formation of  free carbon  (1964) 827 - 831.  in product phase.  Two exothermic DTA peaks  appeared at PO2 kPa for ZrC and in the range of 40 - 1 kPa for HfC. It  = 1  is explained that  the lower  temperature DTA peak is  [4] P. Barnier, F. Thevenot, Eur.  J. Solid State Inorg. Chem. 25  (1988) 495 - 508.  [5] S. Shimada, T.  Ishii,  J. Am. Ceram. Soc. 73 (1990) 2804 -  2808.  [6] S. Shimada, M.  Inagaki, K. Matsui,  J. Am. Ceram. Soc. 75  due to the oxidation of Zr or Hf component of ZrC or  (1992) 2671 - 2678.  HfC with formation of carbon and ZrO2 or HfO2 and  the higher  temperature DTA peak is associated with  the CO2 evolution by oxidation of  free carbon. This  assumption  is  also  supported  from thermodynamic  calculations.  Acknowledgements  The author wishes to express his hearty thanks to  Dr. M. Johnsson at University of Stockholm, Sweden,  for  thermodynamic calculations for oxidation of ZrC  and HfC.  [7] S. Shimada, M. Nishisako, M. Inagaki, J. Am. Ceram. Soc. 78  (1995) 41 - 48.  [8] S. Shimada, M.  Inagaki, M. Suzuki, J. Mater. Res. 11 (1996)  2594 - 2597.  [9] S. Shimada, K. Nakajima, M. Inagaki, J. Am. Ceram. Soc. 80  (1997) 1749 - 1756.  [10] S. Shimada, F. Yunazar, S. Otani,  J. Am. Ceram. Soc. 83  (2000) 721 - 728.  [11] S. Shimada, J. Ceram. Soc. Jpn. 109 (2001) S33 - S42(Special  Review).  [12] D. Gozzi, G. Cascino, S. Loreti, C. Minarini, S. Shimada, J.  Electrochem. Soc. 148 (2001) J15 - J24.  [13] S. Shimada, Oxid. Met. 42 (1994) 357.  [14] D. Gozzi, G. Guzzardi, A. Salleo, Solid State Ionics 83 (1996)  177 - 189.  \\x0c']"
},{
  "_id": 6,
  "PDF": "Ablation behavior and mechanism of C-ZrC, C-ZrC–SiC and C-SiC composites fabricated by polymer infiltration and pyrolysis process.pdf",
  "Text": "['Corrosion Science 86 (2014) 131-141  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Ablation behavior and mechanism of C/ZrC, C/ZrC-SiC and C/SiC composites fabricated by polymer infiltration and pyrolysis process  ⇑  Chunlei Yan, Rongjun Liu  , Yingbin Cao, Changrui Zhang, Deke Zhang  Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 25 February 2014 Accepted 5 May 2014 Available online 13 May 2014  Keywords:  A. Ceramic matrix composites B. SEM C. High temperature corrosion C. Oxidation  1. Introduction  The design and fabrication of materials suitable to withstand ultra-high temperatures and ablative conditions are stimulated by the increasing interest in hypersonic vehicles. One potential materials solution is carbon fiber reinforced composites with matrices containing ultra-high temperature ceramics. In this study, C/ZrC, C/SiC and C/ZrC-SiC composites with different SiC content have been prepared by polymer infiltration and pyrolysis process to investigate potential of ZrC-based composites for ultra-high temperature applications. A comparative study of their ablation behavior and mechanisms is conducted by oxyacetylene torch test. Moreover, a novel insight of self-healing is developed to understand the ablation behavior of C/ZrC-SiC composites.  Ó 2014 Elsevier Ltd. All rights reserved.  External thermal protection systems (TPS) of hypersonic flight vehicles in the form of sharp aerosurfaces such as sharp leading edges and nose caps must be able to withstand high temperature (beyond 2000 °C), high heat flux, severe thermal shock and mechanical stress. So far, there are few, if any, off-the-shelf materials that can meet these future hypersonic thermal protection system needs [1,2]. For example, the conventional refractory alloys, graphite, C/C or C/SiC composites can hardly endure the extreme oxidation and ablation environments for long time, thus novel high temperature materials are required to meet the needs of system (\\x183000 °C). One potential operating at ultra-high temperatures material solution is carbon fiber reinforced composites with matrices containing one or more ultra-high temperature ceramics (UHTCs). UHTCs, which are referred to those refractory transition metal borides, carbides and nitrides with high melting points of over 3000 °C, are of particular interest for ultra-high temperature applications [3]. The initial selection of UHTC materials was based on their melting temperatures, however, the melting points of their oxides are, in fact, more critical. Especially, Zr-based ceramics such as ZrB2 and ZrC have attracted much attention because of their lower cost, broader availability and lower density over other UHTCs. Their high melting points coupled with the ability to form refractory ZrO2 scales give them the capability to withstand temperatures in the 1900-2500 °C range. However, the use of  ⇑ Corresponding author. Tel.: +86 731 84573169; fax: +86 731 84576433. E-mail address: rongjunliu@nudt.edu.cn (R. Liu).  http://dx.doi.org/10.1016/j.corsci.2014.05.005 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.  monolithic components without reinforcements has limitations due to the low fracture toughness and poor thermal shock resistance [1,4]. Continuous fiber reinforced UHTC composites such as C/ZrC-SiC overcome the inherent brittleness of monolithic UHTCs, and can be formed into complex shapes with large size [5,6]. Continuous fiber reinforced ceramic matrix composites can be processed according to: (i) CVI: chemical vapor infiltration; (ii) PIP: polymer infiltration and pyrolysis; (iii) RMI: reactive melt infiltration; (iv) SI-HP: slurry infiltration and hot processing and (v) some hybrid processes in order to associate the advantages of each previous process and to avoid their drawbacks [7-9]. The PIP process involving the use of liquid UHTC precursor is a promising approach, as this technique allows improving homogeneous distribution of UHTCs in the composites and increasing the volume content of UHTCs. In addition, this method can yield near net shape structures which can be of large size and complex shape [10,11]. Considering ceramic matrix composite (CMC) used in oxidizing atmosphere, a significant insight into improving its performance is to utilize the self-healing interphase and matrix. Self-healing materials are able to partially or completely heal damage inflicted on them, and would be ideal for applications which are prone to damage, such as anti-ablation and oxidation fields [12,13]. The concept of self-healing ceramic materials is based on the introduction of elements inside the materials to form fluid oxide phases in a wide range of temperatures that can fill the cracks, protect the fibers in CMC, finally rendering the whole materials self-healing. To date, self-healing of the ceramic materials is mainly achieved by introducing boron species efficient at relatively low temperature (500-1000 °C) and silica-rich phases efficient at higher  \\x0c', '132  C. Yan et al. / Corrosion Science 86 (2014) 131-141  temperature (1000-1500 °C) by forming B2O3 and SiO2 liquid phase, respectively [14,15]. In order to achieve self-healing of the ceramic materials efficient at much higher temperature (>1500 °C), new phases such as Zrand Hf-based additives that can form higher temperature fluid oxide phases need to be introduced. However, no results have been reported on the self-healing mechanism based on the Zrand Hf-based fluid oxide phases at higher temperature (>1500 °C). There are a number of reports in the literature describing the preparation of fiber reinforced UHTC composites for ultra-high temperature applications [5,6,11,16-18]. Paul et al. [16] prepared a range of C/C-UHTC composites by a slurry impregnation and carbon CVI route. Analysis showed that carbon fiber preforms with UHTC powders significantly improved the high temperature oxidation resistance of the composites compared to C-C composites. Hf-based UHTC powders offered superior oxidation protection compared to Zr-based compositions because of the higher melting temperature of HfO2 compared to that of ZrO2. Tang et al. [17] introduced large numbers of ZrB2-based particles into C/C composites to fabricate the C/C-UHTC composites by powder infiltration and isothermal CVI. They reported the effects of UHTC additive, heat flux and ablation time on the ablation behavior of the C/C-UHTC composites. These results highlight the improvement of ablation resistance of C/C composites by introducing UHTCs, while few reports concern the conventional C/SiC composite. Moreover, the UHTC content decreases from external to interior in the C/UHTC composites prepared by slurry impregnation, which is unfavorable for improving the ablation resistance of the composites. Recently, Feng et al. [18] and Zhao et al. [6] reported the ablation behavior of C/C-SiC-ZrC and C/ZrC composites prepared by PIP process respectively, which avoided the drawbacks of slurry impregnation. Cui et al. [19] investigated the microstructure and ablation mechanism of C/C-SiC composites prepared by molten infiltration of silicon powders. It has been noted that while many separate studies were conducted for ablation behavior of the individual ZrCor SiCbased composites [6,11,18,19], the comparative study of C/ZrC, C/SiC and C/ZrC-SiC with different SiC content under the same testing conditions is not necessarily seen. It is also difficult to compare these separate ablation results from one study to another due to changing and some unspecified parameters. Thus, the potential of ZrC-based composites and a shortage of convincing ablation results necessitate the systematic comparative study of the ablation behavior of the C/ZrC, C/SiC and C/ZrC-SiC with different SiC content. The present study is designed to investigate the potential of ZrC-based composites for ultra-high temperature applications. C/ZrC, C/SiC and a range of C/ZrC-SiC composites with different SiC content were prepared by PIP process and their ablation behavior and mechanisms were systematically investigated and compared using an oxyacetylene torch test under the same testing conditions. Besides, a novel insight of self-healing was developed to understand the ablation behavior of the C/ZrC-SiC composites.  [20]. Polycarbosilane (PCS, with a ceramic yield of \\x1865 wt.%) dissolved in xylene with a mass ratio of 1:1 was used as precursor for SiC. For all the composites, the total cycles of PIP process were 20. Firstly, the green composite was obtained by 13 cycles of PIP process with ZrC precursor, the pyrolysis process for the ZrC precursor was conducted at 1000 °C for 1 h under argon atmosphere. Then this composite was divided into 4 parts, and one of the parts was further densified by PIP process with ZrC precursor until 20 cycles to obtain C/ZrC composite. Next, the remaining three parts further densified by 4, 2 and 0 PIP cycles of ZrC precursor separately were used as green products for C/ZrC-SiC composites. All the green composites were heat treated at 1550 °C for 2 h after 13th and the final PIP cycles of ZrC precursor for converting the close pores into open ones and completing the carbothermal reduction to form C/ZrC composites. Finally, the obtained three parts of C/ZrC composites were further densified by 3, 5 and 7 PIP cycles of PCS to obtain C/ZrC-SiC composites with different SiC content and were named as S3-C/ZrC-SiC, S5-C/ZrC-SiC and S7-C/ZrC-SiC, respectively. The C/SiC composite was obtained by 20 cycles of PCS using PIP process, and the pyrolysis of PCS was carried out at 1200 °C for 1 h in flowing argon. The high temperature oxidation testing was carried out in an oxyacetylene torch environment, with approximately 4200 kW/ m2 heat flux and \\x183000 °C flame temperature for the steady flame. The distance between the nozzle tip and the surface of the specimen was 10 mm and the inner diameter of the nozzle tip was 2.0 mm. During the test, the specimens with a size of 30 mm \\x02 60 mm \\x02 10 mm were vertically exposed to the flame for the specified time after the flame was steady. The linear and mass ablation rates were calculated according to the following equations.  Rl ¼ l0 \\x00 l1 t  Rm ¼ m0 \\x00 m1 S \\x01 t  ðaÞ  ðbÞ  where Rl is the linear ablation rate; l0 and l1 are the thickness at ablation center before and after ablation, respectively; t is the ablation time; Rm is the mass ablation rate; m0 and m1 are the weight of sample before and after ablation, respectively; S is the ablated surface area; the ablation rates of the composite were the average value of three samples. After ablation, a few slender bars were cut from the corresponding ablated samples across the ablated center regions to study the evolution of surface and cross-section morphology as well as composition changes of the surface oxide layer. The apparent density of the composites was measured by Archimedes’ method. The phase compositions of the composites were characterized by X-ray diffraction (XRD, D8 Advance, Bruker/Axs Corp., Germany) with Cu Ka radiation. The microstructures of the composites were studied by scanning electron microscopy (SEM, S4800 Hitachi, Japan), equipped with energy dispersion spectroscopy (EDS).  2. Experimental procedure  3. Results and discussion  Polyacrylonitrile-based carbon fibers (T300, 3K, Toray, Tokyo, Japan) were used as the reinforcement and carbon fiber bundles were woven into 3D (3-dimensional, 4-directional) preforms with a \\x1845% fiber volume fraction. Then, in order to obtain fiber protection, a SiC layer was deposited on the fiber preforms at 1100 °C for 2 h by CVD, using methyltrichlorosilane (MTS, CH3SiCl3)/H2 precursor. The liquid ZrC precursor was made from zirconium-containing complex (acetylacetone as the ligand) combined with phenolic resin and had a ceramic yield of \\x1838 wt.%. The details of the preparing procedure for ZrC precursor were described elsewhere  3.1. Property and microstructure of the prepared composites  The samples prepared by PIP process and their properties are summarized in Table 1. The C/SiC composite has a density of 1.95 ± 0.02 g/cm3, while all other composites have higher density in comparison with C/SiC composite, which is attributed to the much higher density of ZrC matrix compared with that of SiC. Besides, the C/ZrC-SiC composites with more PIP cycles of PCS have a higher density because of the higher ceramic yield and impregnation efficiency of PCS despite the same total PIP cycles.  \\x0c', 'C. Yan et al. / Corrosion Science 86 (2014) 131-141  133  Table 1 Samples prepared by PIP process and their properties.  Samples  Fiber fraction (vol%)  Density (g/cm3)  Open porosity (%)  ZrC fraction (vol%)  SiC fraction (vol%)  C/ZrC S3-C/ZrC-SiC S5-C/ZrC-SiC S7-C/ZrC-SiC C/SiC  45.7 45.7 45.7 45.7 46.0  2.01 ± 0.02 29 ± 4 2.04 ± 0.04 22 ± 3 2.07 ± 0.03 18 ± 2 2.13 ± 0.02 15 ± 2 1.95 ± 0.02 9 ± 1  22.3 ± 0.5 18.9 ± 0.4 16.6 ± 0.2 14.1 ± 0.2 -  - 10.8 ± 0.1 16.9 ± 0.2 22.9 ± 0.2 42.6 ± 0.6  Based on the aforementioned reason, the C/ZrC-SiC composites with more PIP cycles of PCS have a lower open porosity, and their SiC volume fractions increase rapidly accordingly with little change of ZrC content. As shown in Figs. 1 and 2, it can be observed from the diffraction patterns before ablation that the ZrC is the only crystal phase in C/ ZrC composites, while ZrC and SiC are the main and minor crystal phase in C/ZrC-SiC composite, suggesting that the carbothermal reduction was completed and the ZrC matrix formed. At the pyrolysis and high temperature heat treatment stages, the evaporation of solvent and release of gaseous products result in a significant loss of weight of the precursor. Meanwhile, the transformation of the liquid ZrC precursor having a typical density of \\x181 g/ cm3 into crystal ZrC ceramic products of higher density (\\x185 g/cm3) after the high temperature treatment implies volume shrinkage of the matrix with production of large numbers of pores. The microstructure of C/ZrC composite is characterized by many interbundle cracks (Fig. 3a) and intrabundle micropores (Fig. 3b), which have a negative impact on the mechanical as well as the ablation property. So the composites were further densified by PIP process of PCS. Although cracks still exist in the interbundle areas for S7-C/ZrC-SiC composite after 7 PIP cycles of PCS compared with the dense matrix of C/SiC composite, the densification of S7-C/ZrC-SiC is remarkably improved in comparison with C/ZrC composite (Fig. 4). Despite large numbers of pores existing in the composite, the ZrC phase in C/ZrC composite obtained by PIP process disperses homogeneously in the matrix (Fig. 3), which is important to ensure a uniform and lasting protection against ablation during testing. Meanwhile, the volume content of the ZrC phase is improved, as the liquid ZrC precursor is easily inserted into both interbundle and intrabundle zones.  3.2. Ablation property and macromorphology  The high temperature performance of all the obtained composites was studied utilizing a custom built oxyacetylene torch test rig. Table 2 summarizes the torch temperature, mass ablation rate  Fig. 1. XRD patterns of C/ZrC composite before and after ablation.  Fig. 2. XRD patterns of S7-C/ZrC-SiC composite before and after ablation.  and linear ablation rate after 80 s oxyacetylene torch tests. It can be concluded that both of the mass and linear ablation rates are lowest for the C/ZrC composite, while the introduction of SiC degrades the ablation property of C/ZrC composite. After ablation for 80 s, the linear and mass ablation rates for C/ZrC composite are 0.013 ± 0.001 mm/s and 0.13 ± 0.01 mg cm\\x002 s\\x001, respectively. With the increase of SiC content, the linear and mass ablation rates increase accordingly. S7-C/ZrC-SiC composite with a higher SiC content has a linear ablation rate of 0.026 ± 0.002 mm/s and a mass ablation rate of 0.69 ± 0.04 mg cm\\x002 s\\x001, which are still superior to that of C/SiC composite. The C/SiC composite shows the poorest ablation property among all the composites, and its linear and mass ablation rates are 0.061 ± 0.003 mm/s and 1.46±0.05 mg cm\\x002 s\\x001, which are nearly 5-fold and 11-fold of that of C/ZrC composite. Fig. 5 shows the mass loss and linear recession rates vs. time for S3-C/ZrC-SiC and C/SiC composites. The time has little effect on the mass loss rate of C/SiC, as the ablation progresses uniformly without any anti-ablation layers formed. However, the mass loss rate of S3-C/ZrC-SiC composite after 80 s ablation is twice of that after 40 s. At the early ablation stage of S3-C/ZrC- SiC composite, on the one hand, the oxidation of the fibers and evaporation of low melting point products result in mass loss, on the other hand, the oxidation of ZrC matrix inducing the forming of ZrO2 leads to mass increase. As a result, the mass loss rate after 40 s ablation is much lower than that after 80 s, as the ablation in later period only leads to mass loss after the completion of the main formation stage of ZrO2. However, it can be observed from Fig. 5b that the time has the same effect on the linear recession rates of C/SiC and S3-C/ZrC-SiC composites. Both of the linear recession rates slightly decrease with increase of time, as the ablation temperature decreases for the longer distance between the flame tip and ablated surface with the progress of the ablation. Photographic images of the obtained composites after 80 s oxyacetylene torch testing are shown in Fig. 6. Oxyacetylene torch testing is a very aggressive test because it involves high temperature, high-velocity gas flow and severely oxidizing environment. The flame with a temperature of \\x183000 °C is vertically exposed to the sample, and heating rates of up to approximately 500 °C s\\x001 can be achieved, moreover, the initial cooling rate is \\x181000 °C s\\x001 after extinguishing the flame [21]. Nevertheless, all the composites survived the high heating/cooling rates and maintained an integral structure, indicating excellent thermal shock resistance for all composites. The reaction products formed after the oxyacetylene torch testing are characterized using XRD as shown in Figs. 1 and 2. The XRD results reveal that the surface oxide layer only contains monoclinic zirconia (m-ZrO2) for both C/ZrC and S7-C/ZrC-SiC composites. During the ablation procedure, the width of torch is 4-5 mm, and a sharp temperature gradient  \\x0c', '134  C. Yan et al. / Corrosion Science 86 (2014) 131-141  Fig. 3. SEM micrographs on the polished cross section of C/ZrC composite: (a) interbundle areas and (b) intrabundle areas.  Fig. 4. Typical SEM micrographs on the polished cross section of (a) S7-C/ZrC-SiC and (b) C/SiC composites.  Table 2 Summary of the results for C/ZrC, C/ZrC-SiC and C/SiC composites after oxyacetylene torch testing.  Sample  Torch temperature (°C)  C/ZrC S3-C/ZrC-SiC S5-C/ZrC-SiC S7-C/ZrC-SiC C/SiC  \\x183000 \\x183000 \\x183000 \\x183000 \\x183000  Ablation time (s)  Mass ablation rate (mg cm\\x002 s\\x001)  Linear ablation rate (mm/s)  80 80 80 80 80  0.13 ± 0.01 0.32 ± 0.02 0.42 ± 0.03 0.69 ± 0.04 1.46 ± 0.05  0.013 ± 0.001 0.016 ± 0.001 0.021 ± 0.001 0.026 ± 0.002 0.061 ± 0.003  from the center region to outer parts takes place. The presence of molten ZrO2 phases (with a melting point of 2715 °C [16]) in the center region can be seen on all of the ZrC-based composites, independently confirming the temperature of the flame. As far as the C/ ZrC composite is concerned, it can be observed that the entire surface is oxidized and could be divided into two regions (Fig. 6a), i.e. the center molten region (area a in Fig. 7a) with a diameter of \\x185 mm, the outer oxidation region (area b in Fig. 7a). As discussed elsewhere [22], the ZrC starts oxidation at about 300 °C, so in a high temperature and oxidizing environment, the ZrC matrix on the whole surface oxidizes to form an ZrO2 layer. Besides, this ZrO2 layer is little adherent and falls off easily; possibly because of the absence of any glassy phases during the test and the build up of CO/CO2 gas at the interface of substrate-oxide layer. However, the photograph of the cross section along the ablated center of C/ZrC composites shows that the interior of composites is not oxidized severely (Fig. 6b), as the oxygen barely diffuses into the interior zone for the blocking effect of the surface oxide layer. As shown in Fig. 6c, a big pit (denoted as zone 1 in Fig. 7b) with a diameter of \\x1825 mm and a depth of \\x185 mm without covering of any oxide layer appears in the center of the C/SiC composite, revealing the severe ablation of C/SiC composites. A circle of glass-like layer covers the edge of the ablated pit despite falling off of some parts (denoted as zone 2 in Fig. 7b); while the outer region with covering of frost-like oxide product remains nearly untouched. Therefore, three typical regions as shown in Fig. 7b  correspond to the morphology of the ablated surface of C/SiC composite. It can be observed from Fig. 6d-f that the morphology of C/ZrC-SiC composites is quite similar, and a schematic drawing of the ablated surface abstracted from the photographs of all C/ZrC-SiC composites is shown in Fig. 7c. The surface morphology of C/ZrC-SiC composites can be divided into four regions from center to external, i.e. center molten region (zone A), white oxide region (zone B), gray edge region (area C), and the outermost nearly untouched region (zone D), which associates the division of morphology for both C/SiC and C/ZrC (Fig. 7). It is noteworthy that with increase of SiC content, the center molten zone is enlarged, which may be attributed to the lower melting temperature of the oxide layer resulting from the solution of SiO2 in ZrO2 [23]. Moreover, the outermost region of the ablated C/ZrC-SiC composites (D zone in Fig. 7c) represents less oxidized with increase of SiC content compared with that of C/ZrC composites, indicating the improvement of anti-oxidation property by introducing SiC. Finally, a common interesting feature can be summarized as follows from the morphology of ZrC-based composites. For all these composites, the center region is covered by molten ZrO2 layers without any bare composite despite the existence of scouring from the combustion flow, indicating a good protection achieved by the molten ZrO2 layer.  3.3. Microstructure  The differences in ablation properties and macromorphology imply different microstructure of the ablated composites. In order to study more details of the ablation behavior for the obtained composites, surface and cross-section microstructures of the composites after ablation are investigated and discussed. The analysis of surface microstructure is based on macromorphology and the division of the ablated surface of the obtained composites as shown in Fig. 7, and the S7-C/ZrC-SiC sample among C/ZrC-SiC composites is utilized as the typical sample for microstructure analysis.  \\x0c', 'C. Yan et al. / Corrosion Science 86 (2014) 131-141  135  Fig. 5. Mass loss rates (a) and linear recession rates (b) of C/SiC and S3-C/ZrC-SiC composites vs. time.  Fig. 6. Photographs of the obtained composites after 80 s oxyacetylene torch testing: (a) C/ZrC, (b) cross section of (a) along the ablated center, (c) C/SiC, (d) S3-C/ZrC-SiC, (e) S5-C/ZrC-SiC and (f) S7-C/ZrC-SiC. The length and width of the composites are 60 mm and 30 mm.  Fig. 7. Schematic drawings of composites).  the ablated surface of  the obtained composites (the drawing of C/ZrC-SiC is abstracted from the photographs of S3-, S5-, S7-C/ZrC-SiC  3.3.1. C/ZrC composite  The microstructures of the ablated surface for C/ZrC composites are shown in Fig. 8. As shown in Fig. 8a, some oxide layers covering the fiber bundles and inter bundle areas can be observed despite the spallation of the most surface molten ZrO2 layers in the sample preparation for SEM analysis. The inset in Fig. 8a shows the  microstructure of the stripped molten ZrO2 layer, which is a continuous layer without any pore distribution and endowed with glass feature for its flowing and smooth appearance. It could be anticipated that an oxygen diffusion barrier and an excellent protection against ablation can be achieved by these integral molten ZrO2 layers observed from as-received sample after testing shown  \\x0c', '136  C. Yan et al. / Corrosion Science 86 (2014) 131-141  Fig. 8. SEM morphologies of the surface areas for the ablated C/ZrC composite: (a) center molten region; (b) outer oxidation region and (c) carbon fibers underneath the surface oxide layers in center region. The inset shows the microstructure of the stripped molten ZrO2 layer.  in Fig. 6a. The microstructure of the oxide layer on the outer oxidation region is shown in Fig. 8b. The ablated surface in this region is covered by ZrO2 grains which have obvious boundaries among each other and signs of necking, indicating a lower ablation temperature in this region. This oxide layer is porous, which is attributed to little sintered degree and the escape of the CO/CO2 gas produced by oxidation reaction. Fig. 8c shows the coincidence of the needle-like fibers and the white oxide layer which wraps fibers, indicating the severe oxidation of both fibers and matrix in center molten region. Fig. 9 shows the morphologies of the cross section for center molten region of the ablated C/ZrC composite. The cross section can be divided into three different regions from surface to interior, i.e. matrix-depleted layer, transition layer and the untouched composite (Fig. 9a). In the matrix-depleted layer, the matrix in fiber bundles is depleted by oxidation, thus the partially oxidized fibers remain; however, as shown in Fig. 9b, the crack between the fiber bundles is full of oxide products. Firstly, the matrix migrates to the surface to form the molten oxide layer when the composite is exposed to the flame. Secondly, some of the matrix migrates to the interbundle areas to be oxidized, as the crack between the fiber bundles provides the oxygen diffusion channel. The oxide layer  formed both on the surface and in the interbundle areas can act as a barrier for ablation and further diffusion of oxygen. The transition region is characterized by partially depleted matrix, as the oxygen seldom infiltrates into this region. Then the next layer is the untouched composites.  3.3.2. C/SiC composite  The microstructures of the surface areas for the ablated C/SiC composite are shown in Fig. 10. In the ablated pit areas, only the bare fiber bundles and the exposed pores can be observed, and there are no oxide layers covering the surface of the pit (Fig. 10a). As shown in Fig. 10b, the fibers are severely oxidized into needle shape, and there is no matrix remained among fibers in addition to the interbundle areas, as the flame can hardly reach theses areas for the sheltering effect of the bundles. The matrix in the interbundle areas has a remarkable sign of erosion by the oxyacetylene flame; moreover, microcracks take place due to the severe thermal shock during testing. The edge areas around the ablated pit are covered by a viscous silica-containing glass layer, which is confirmed by the inset EDS pattern (Fig. 10c). It can be observed from Fig. 10c that the fibers underneath the SiO2 layer remain less oxidized, as the oxygen can hardly infiltrate this glassy  Fig. 9. SEM morphologies of (a) cross section and (b) magnified image of the corresponding areas in (a) for center molten region of the ablated C/ZrC composite.  Fig. 10. SEM morphologies of the surface areas for the ablated C/SiC composite: (a and b) the ablated pit areas in the center region; (c) the edge areas around the ablated pit. The inset image shows the EDS pattern of the glassy layer.  \\x0c', 'C. Yan et al. / Corrosion Science 86 (2014) 131-141  137  Fig. 11. SEM morphologies of (a) cross section and (b) magnified image of fiber bundles in (a) for center region of the ablated C/SiC composite.  barrier. The oxide products concerning the oxidation of SiC matrix are SiO2 (s or l) for the conventional oxidation reaction and SiO (g) for the active reaction. The SiO2 (l) formed in the center region evaporates and recondenses on the edge of the pit. Meanwhile, SiO (g) recombines with oxygen and condenses into a silica-rich glassy layer for lower temperature experienced in this region. Besides, the native oxidation of SiC matrix also leads to formation of SiO2 (l). As a result, a protective silica-rich layer forms around the edge of the pit. Fig. 11 shows the microstructures of the cross section for center region of the ablated C/SiC composite. There should have been bare fibers on the top edge of the cross section for ablated C/SiC composite as expected from SEM morphologies of the surface areas shown in Fig. 10. However, these bare fibers without bonding of matrix resulting from the oxidation are easily broken, and then fall off when prepared the sample for SEM study. As confirmed by the enlarged image of fiber bundles in cross section shown in Fig. 11b, there are no bare fibers left. As shown in Fig. 11b, there are two layers, i.e. the matrix partially depleted layer and untouched composite, underneath the surface bare fiber layers. No visible boundaries can be observed between these two layers because a uniform temperature gradient may appear from surface to interior of composite.  3.3.3. C/ZrC-SiC composite  The typical morphologies and the corresponding EDS analysis of the oxide layer on the ablated surface for S7-C/ZrC-SiC composite are shown in Fig. 12. As shown in Fig. 12a, a continuous molten layer containing only ZrO2 (confirmed by EDS) appears on the surface of the center region with sporadic distribution of micropores. Moreover, there are no microcracks in this ZrO2 layer. However, most of the molten ZrO2 fell off during preparing sample for SEM analysis. Considering the photographs of ablated S7-C/ZrC-SiC composite shown in Fig. 6f, the as-received integral glassy ZrO2 structures resulting from melt of ZrO2, which flow and cover the surface, seal the pores and cracks and block the diffusion channels of oxygen, give the composite a good anti-ablation property during testing. Carbon fibers underneath the surface molten oxide layers in center region present needle-shape (Fig. 13a), indicating a severe oxidation. However, unlike the ablated C/SiC composite, a quite large amount of matrix remained among fibers, which is beneficial to the supplement of the surface oxide layers. Fig. 12b shows the micrographs of oxide layer in the outer white oxide region, which is composed of Zr, O, and a trace of Si confirmed by the corresponding EDS pattern. With a longer distance away from flame tip, the temperature experienced in this region is lower, which results in a small amount of remnant SiO2 in addition to the major ZrO2 phase. As all we know, SiC undergoes active oxidation at reduced oxygen partial pressures above \\x181500 °C leading to the gaseous SiO and CO without forming a SiO2 scale [24]. In addition, the vapor pressure of SiO2 at the testing  temperature is much higher than that of ZrO2 [23,25]. Hence the SiO2 is not stable enough to remain in the oxide layer and cannot offer any additional protection in the center and outer white oxide regions. However, the remnant SiO2 lowers the eutectic temperature of the oxide scale, which leads the oxide particles to have the capability to be sintered at a relatively lower temperature. It can be observed from the inset enlarged image shown in Fig. 12b that good bonding between the particles and signs of fusing take place, which lead to well-sintered oxide scale against ablation and oxidation. It could be anticipated that with much higher temperature experienced, this sintered scale will be melted to be a continuous scale to provide further protection for the composite. The carbon fibers and matrix underneath the oxide layer in this region present less oxidized (Fig. 13b) compared with that in center region, which indicates that a good protection by the oxide scale has been achieved. Fig. 12c shows the microstructure and corresponding EDS pattern in gray edge region. This gray oxide layer is composed of silica confirmed by the EDS pattern, which is amorphous phase because no peaks of SiO2 in the XRD patterns are observed, indicating a mild ablation environment in this region. This silica-containing glass layer is well-adherent to the substrate and can act as a blocking barrier for the oxygen diffusion. The distribution of pores on this oxide layer is believed to be due to the escape of SiO and CO/CO2 gases generated as a result of the oxidation of ZrC and SiC. It is noteworthy that the molten silicacontaining scale also exists below the surface ZrO2 layer (the arrows denote in Fig. 12c) because of a lower temperature experienced under the ZrO2 layer. This molten silica-containing phase can act as a healing agent under the ZrO2 scale to seal the pores and cracks, which is beneficial to the improvement of the oxidation resistance. The micrographs of the cross section for center molten region of the ablated S7-C/ZrC-SiC composite are shown in Fig. 14. It can be concluded from Fig. 14 that the cross section can be divided into three regions from external to interior, i.e. the molten ZrO2 layer, the matrix depleted region and the unaltered C/ZrC-SiC composite. The thick molten ZrO2 scale spreads on the top edge of the cross section and is regarded as an ultimate layer against ablation. Compared with C/ZrC composite, the matrix depleted region of S7-C/ ZrC-SiC composite is thinner (Fig. 14b), which is attributed to good anti-oxidation resistance of SiC matrix as well as the sharp temperature gradient from surface into interior. The weak bonding of the matrix can be observed in the untouched composite, which may be caused by the sharp thermal shock during the high temperature testing.  3.4. Ablation mechanism  The details concerning the ablation property, macromorphology, and microstructure give much original material for understanding the ablation behavior and its corresponding mechanisms of the  \\x0c', '138  C. Yan et al. / Corrosion Science 86 (2014) 131-141  Fig. 12. SEM morphologies and EDS analysis of the surface oxide layers for the ablated S7-C/ZrC-SiC composite: (a) center molten region; (b) white oxide region and (c) gray edge region corresponding to zone 1, 2, 3 in Fig. 6f, respectively.  Fig. 13. Ablation morphologies of carbon fibers underneath the surface oxide layers for S7-C/ZrC-SiC composite: (a) center molten region and (b) white oxide region.  Fig. 14. SEM morphologies of (a) cross section and (b) magnified image of the corresponding areas in (a) for center molten region of the ablated S7-C/ZrC-SiC composite.  \\x0c', 'C. Yan et al. / Corrosion Science 86 (2014) 131-141  139  obtained composites, which is beneficial to the selection of composites used in ultra-high temperature environment. The reactions schemes for the oxidation of the various composites are given below [6,8,9,16-18]:  ZrCðsÞ þ 2O2 ðgÞ ! ZrO2 ðsÞ þ CO2 ðgÞ  ZrCðsÞ þ 3 2  SiCðsÞ þ 3 2  O2 ðgÞ ! ZrO2 ðsÞ þ COðgÞ  O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ  SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  SiO2 ðlÞ ! SiO2 ðgÞ  CðsÞ þ O2 ðgÞ ! CO2 ðgÞ  2CðsÞ þ O2 ðgÞ ! 2COðgÞ  ð1Þ  ð2Þ  ð3Þ  ð4Þ  ð5Þ  ð6Þ  ð7Þ  The diagram of the ablated C/ZrC composite summarized from both macromorphology and microstructure of cross section is shown in Fig. 15. The cross section of the ablated C/ZrC composite is divided into 4 layers and the uppermost layer is the thick molten ZrO2 layer, while the surface is composed of 2 regions, i.e. the center ablated pit covered by molten ZrO2 and the outer porous ZrO2 layer. The oxidation of the composite involves the oxidation of the carbon fibers and ZrC matrix (reactions (1), (2), (6), (7)), and the oxidation of the carbon fibers contributes mainly to the mass loss of the composite. On the other hand, the erosion in the center region of the composite is the result of multi-factors. The stripping of the molten ZrO2 because of scouring from the high-velocity combustion flow, the volume shrinkage of the ZrO2 layer resulting from the melting and sintering, and subsidence of the ZrO2 layer because of large porosity of the composite, as well as the pressure of combustion gas imposed on the ZrO2 layer account for the erosion of the ablated composite together. The C/SiC composite shows a different ablation behavior as shown in Fig. 16 in comparison with C/ZrC. The C/SiC composite shows the severest ablation performance with the largest ablated pit and deepest erosion among other composites. SiC undergoes active oxidation above \\x181500 °C leading to the formation of gaseous SiO and CO [24]. In addition, the vapor pressure of SiO2 at the testing temperature is very high [25], which leads the liquid SiO2 to gasify to form a gaseous products [26]. Moreover, the oxidation of carbon fibers results in the gaseous CO and CO2. As a result, all the gaseous products resulting from oxidation and high temperature lead to the both high mass loss rate and linear recession rate. Although the glass SiO2 scale forms on the edge of the ablated pit where the sample experienced <2000 °C, no stable refractory oxide layers cover the surface of the  Fig. 16. Diagram of the ablated C/SiC composite.  ablated pit. Thus the more new fibers and matrix are exposed to the flame to be gasified, which lead the ablation of the C/SiC composite into a vicious circle. However, compared with C/ZrC composite, the C/SiC composite beyond the ablated pit remains less oxidized. Considering the above results and discussion concerning the ablation property and morphology, the ablation behavior of C/ZrC-SiC composite combines the features of both C/SiC and C/ ZrC composites, and its ablation diagram is shown in Fig. 17. The C/ZrC-SiC composite with lower content of SiC (S3-C/ZrC-SiC) has a similar macromorphology to that of C/ZrC, and both of them show a small ablated pit and severely oxidized surface (Fig. 6a and d). On the other hand, the C/ZrC-SiC composite with higher content of SiC (S7-C/ZrC-SiC) presents a similar macromorphology to that of C/SiC showing a silica-containing circular layer and outer untouched composite (Fig. 6c and f). The size of the center region and outer ZrO2 region for C/ZrC-SiC composite is nearly equal to that of the ablated pit for C/SiC composite (Figs. 16 and 17), indicating that the temperature experienced in these regions is so high that the SiC and its oxidation products are not stable enough with the exception of ZrO2 phase. As a result, the introduction of SiC degrades the ablation property of C/ZrC composite. The oxidation of SiC and carbon fibers leading to the gaseous products contributes mainly to the mass loss of the C/ZrC-SiC composite. The spallation of the molten ZrO2 because of scouring from the highvelocity combustion flow and the depletion of the SiC and carbon fibers due to oxidation account for the erosion in the center region of the C/ZrC-SiC composite. To draw a conclusion, the ablation and oxidation behavior of ZrC and SiC during testing are different. So the roles of ZrC and SiC in the ablation process should be further discussed. The ZrC can be regarded as the source of the refractory ZrO2 phase, which forms stable molten ZrO2 scale against ablation in the center region, where the sample experiences >2700 °C. Nevertheless, C/ZrC composite shows poor oxidation resistance and the oxide  Fig. 15. Diagram of the ablated C/ZrC composite.  Fig. 17. Diagram of the ablated C/ZrC-SiC composite.  \\x0c', '140  C. Yan et al. / Corrosion Science 86 (2014) 131-141  layer is little adherent and falls off easily. Experimentally, addition of SiC can significantly improve the oxidation resistance in the intermediate temperatures that range from 1200 to \\x181700 °C by forming a protective silica layer [27]. The C/ZrC-SiC composites were obtained by introduction of SiC using PIP process, the oxidized areas of C/ZrC-SiC composite with higher content of SiC are limited to the center and outer ZrO2 region (Fig. 17), much smaller than that of C/ZrC, indicating that the SiC can act as an anti-oxidation constituent. Moreover, the glassy silica phases underneath the outer ZrO2 layer and the untouched composite endowed by SiC matrix in C/ZrC-SiC composite can stick the central ZrO2 layer to keep its integrity, which is superior to C/ZrC composite. Therefore, as far as both the anti-oxidation and quite good ablation properties are concerned, the C/ZrC-SiC composite with a certain content of SiC is a better choice. Combining the features of both the ZrC and SiC matrix, a high temperature (>2000 °C) self-healing mechanism has been achieved by C/ZrC-SiC composite. The fluid ZrO2 scale, regarded as healing agent, covers the center region of the sample and seals the cracks and pores to impede the in-depth oxygen diffusion toward the oxidation prone fibers, rendering the whole materials self-healing against ablation and oxidation. The areas beyond the center region of C/ZrC-SiC composite where the sample experiences lower temperature than the center region are characterized by sintered ZrO2 scale and molten glassy SiO2 underneath the surface ZrO2 scale. Besides, the edge around the outer sintered ZrO2 scale is covered by a silica-containing layer. In those areas, the sintered ZrO2 scale and molten glassy SiO2 can act as healing agents to seal the pores and cracks and protect the fibers against oxidation in the intermediate temperature range. The self-healing mechanism can be achieved provided that the continuous supplement of the fluid oxide scale can be obtained to strike a better balance between the depletion and the supply of the healing agents. As all we know, the UHTC content decreases in the C/UHTC composites prepared by slurry impregnation from external to interior because of the increasing blocking effect of the fibers. As a result, an ineffective supplement of fluid refractory oxide takes place for the shortage of UHTC phase in the interior of the composite. Thus the interior bare composite is exposed [16,17], as the depletion of surface oxide takes place due to the spallation and evaporation of oxidation products with the ablation proceeding. However, the supplement of ZrO2 in the C/ZrC-SiC composite by PIP process is not worth the attention, because the ZrC matrix obtained from liquid ZrC precursor gets a homogeneous distribution in the composite. Therefore, the ZrC matrix in composite acts as a reservoir of the ZrO2 phase, and then supplements the ZrO2 loss continuously in the whole testing, resulting in a continuous self-healing without any bare composite exposed. On the other hand, the ablation of C/SiC goes into a vicious circle because no self-healing mechanism can be obtained for the absence of stable molten oxide layer, while the severe oxidation of C/ZrC composite gives it limitations for further applications in ultra-high temperature environment.  4. Conclusions  The potential of ZrC-based composites for ultra-high temperature applications has been studied in this work by comparing the ablation behavior of C/ZrC, C/SiC and C/ZrC-SiC composites prepared by PIP process. Both the mass and linear ablation rates are lowest for the C/ZrC composite. With the increase of SiC content, the linear and mass ablation rates of C/ZrC-SiC composites increase accordingly, nevertheless, they are still much superior to that of C/SiC composite. The ZrC is regarded as the source of the refractory ZrO2 phase, which forms stable molten ZrO2 scale against ablation in the center region, where the sample experiences >2700 °C. Nevertheless,  C/ZrC composite shows poor oxidation resistance, besides, the oxide layer is little adherent and falls off easily. The SiC and its oxidation products are not stable enough to survive the high temperature in the center region. However, the SiC can act as an anti-oxidation constituent in those areas where samples experience lower temperatures (<1700 °C). Therefore, the oxidized areas of C/ZrC-SiC composite with higher content of SiC are limited to the center and outer ZrO2 region, much smaller than that of C/ZrC. Combining the features of both ZrC and SiC matrix, the C/ZrC-SiC composite with a certain content of SiC is one potential material for ultra-high temperature applications. A high temperature self-healing mechanism has been achieved by C/ZrC-SiC composite in oxidizing environment. The ZrC matrix homogeneously distributed in composite acts as a reservoir of the ZrO2 phase, and then supplements the ZrO2 loss in the center region continuously in the whole testing, resulting in a continuous self-healing without any bare composite exposed.  Acknowledgments  This work was financially supported by the National Natural Science Foundation of China (no. 51102282) and Aid Program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province.  References  [4]  [7]  [1] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for, 2000°C+ hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. 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},{
  "_id": 7,
  "PDF": "Ablation behavior of ZrB2–SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions.pdf",
  "Text": "['Available online at www.sciencedirect.com  Composites Science and Technology 68 (2008) 1718-1726  COMPOSITES SCIENCE AND TECHNOLOGY  www.elsevier.com/locate/compscitech  Ablation behavior of ZrB2-SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions  Xinghong Zhang *, Ping Hu *, Jiecai Han, Songhe Meng  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, PR China  Received 4 June 2007; received in revised form 17 December 2007; accepted 6 February 2008  Available online 15 February 2008  Abstract  ZrB2-20vol%SiC ultra high temperature ceramic (UHTC) was prepared by hot-pressing. Ablation tests of the flat-face models were conducted under ground simulated atmospheric re-entry conditions using arc-jet testing with heat fluxes of 1.7 MW/m2 and 5.4 MW/m2 under subsonic conditions, respectively. There was little weight or configuration change after ablation at a heat flux of 1.7 MW/m2. However, ZrB2-SiC composite underwent severe ablation and whose surface temperatures exceeded 2300 °C at a heat flux of 5.4 MW/m2. Sharp-shape leading edge models were ablated under supersonic conditions with the stagnation pressure and Mach number of  1.2 atm and 2.7 M, respectively, and sharp-shaped leading edge C/SiC models were also ablated under the same condition for compar ison. ZrB2-SiC composite exhibited an excellent compared with C/SiC material. Results indicate that ZrB2-SiC ultra high temperature ceramics are the potential candidates for leading edges. The temperature limit for UHTC is controlled by the softening and degradation of the formed oxide scale.  thermal-oxidative and configurational stability in the simulated re-entry environment  Ó 2008 Elsevier Ltd. All rights reserved.  Keywords: A. ZrB2; A. SiC; B. Ablation behavior  1. Introduction  to endure higher  surface  temperatures  [9-12]. The  lower  drag should improve ascent performance, while the higher  Ultra high temperature ceramics (UHTC) include some  lift to drag produces a greater range for abort and re-entry  refractory metal diborides  that were historically  studied  trajectories.  Improved range provides operational  safety  and developed since the 1960s and early 1970s [1-3]. Inter benefits for the system. Sharp leading edges have not been  est  in UHTC has  increased significantly  in recent years  used for reusable vehicles in the past because materials able  because of the drive to produce thermal protection systems  to repeatedly withstand the higher entry and abort temper (TPS)  and  other  components  for  hypersonic  aerospace  atures were not available. Recent developments in the use  vehicles [4-8].  of UHTC have given renewed hope to the eventual realiza In contrast to traditional blunt capsules or Shuttle-type  tion of sharp leading edge vehicles [13-15].  vehicles,  characterized  by  poor  gliding  capabilities,  the  future use of UHTC opens new prospects for the develop ment of space planes with slender fuselage noses and sharp  wing leading edges. Vehicles with sharp leading edges and/  or  sharp noses have  lower drag and higher  lift  to drag  However, report on the ablation behavior of ZrB2-SiC under simulated atmospheric re-entry conditions is limited  and their surface temperatures are always less than 2000 °C. In this study, ZrB2-20vol%SiC ultra high temperature ceramic was prepared by hot-pressing. The flat-face  ratios (L/D)  than blunt-nosed vehicles, but  they also have  models  together with sharp-shaped leading  edge models  * Corresponding authors. Tel./fax: +86 45186402382.  E-mail addresses:  zhangxh@hit.edu.cn (X.H. Zhang), huping@hit.e du.cn (P. Hu).  0266-3538/$ see front matter Ó 2008 Elsevier Ltd. All rights reserved.  doi:10.1016/j.compscitech.2008.02.009  were tested using an arc-jet  facility under different condi tions, and sharp-shaped leading edge C/SiC models were  also ablated under  the  same  condition for  comparison.  The purpose of  this  study  is  to investigate  the ablation  \\x0c', 'behavior  of  the ZrB2-SiC composite mechanism is also discussed.  and  the  ablation  2. Experimental procedure  The samples used here for oxidation testing were fabri cated  from commercial  ZrB2 Non-ferrous Metal Research, China)  (Northwest  Institute  for  and SiC (Weifang  Kaihua Micro-powder Co., Ltd., China)  powders. The  ZrB2 and SiC powders had the same purity of 99.5%, with mean particle size of 5 lm and 2 lm, respectively. The  powder mixture of ZrB2 + 20vol% SiC was ball-milled in ethanol for 8 h using hard milling media and dried in a  rotating  evaporator. Milled powder was  then uniaxially  hot pressed in a boron nitride coated graphite die at 2000 °C for 60 min under vacuum and 30 MPa of applied  pressure. The heating schedule has been described in detail  elsewhere  [16]. Bulk density and theoretical density were  evaluated using the Archimedes method and the  rule-of mixture, respectively.  ZrB2-SiC models with a cylindrical shape were cut from the billet, and then exposed to sustained enthalpy flows  using an arc-jet  facility with an average total enthalpy in  the order of 10-30 MJ/kg. The facility is equipped with a  60 KW plasma torch which has a diameter of 30 mm of  the exit nozzle that can be operated in air with a stagnation  pressure of 0.03-0.5 atm. The dimension of the model is U20 \\x02 30 for the high heat flux under subsonic conditions and that of U30 \\x02 30 for the low heat flux. The ablation time for these models is 600 s. The heat fluxes were mea sured using flat-faced Gardon-type calorie meter and the  measured fluxes 5.4 MW/m2,  for  two  flat-face models were  1.7  and  respectively. The sharp leading edge models  (Fig. 1) were also cut from the billet and then tested under  supersonic conditions using an arc-heated wind tunnel that  is operated in air. The size of the exit nozzle of 64 mm \\x02 32 mm and supersonic flow was controlled in the temperature range of 1440-1450 °C. The stagnation pressure and Mach num the wind  tunnel  is  the  temperature  of  the  ber of  the supersonic flow at  the exit nozzle were 1.2 atm  and 2.7 M,  respectively. The average  total enthalpy is  in  the order of 2-4 MJ/kg.  The  experiments were  carried out with a  two-colour  Raytek pyrometer  (RAYMR1SCSF, USA) which covers  a  temperature  range of  1000-3000 °C. X-ray diffraction  (Rigaku, Dmax-rb)  and  scanning  electron microscopy/  energy dispersive spectroscopy (FEI Sirion, Holand) were  used to characterize the phase composition and the micro structure of the surface and the cross-section of the samples  after testing.  3. Results and discussion  3.1. Microstructure of  the hot-pressed composite  The bulk density of the sintered ZrB2-SiC billets was 5.41 g/cm3, which corresponds to a relative density higher  than 98%. Fig. 2 shows a scanning electron micrograph  of  the polished surface of  the ZrB2-SiC composite. The microstructure of the composite is regular, in which the mean grain size is about 6 lm and the residual porosity is  very scarce as shown in Fig. 2. The grain size of this mate rial was estimated from the fracture surface of the compos ite (not shown). The SiC particulates, which represent  the  majority of the dark features in Fig. 2, are homogeneously  dispersed in the ZrB2 matrix and no agglomeration was detected. Exaggerated grain growth of ZrB2 was restricted due to the existence of SiC particles. In addition, the intro duction of SiC substantially enhanced densification of ZrB2 during hot-pressing. ZrO2 and B2O3 were assumed as the main oxygen carriers upon the surfaces of ZrB2. Such a contamination by oxygen promotes vapor phase transport  at  temperatures below which mass  transfer mechanisms  like  boundary/volume  diffusion, which  are much more  effective  than vapor phase  for densification,  start  taking  place;  the  anticipated  coarsening  decreases  the  driving  force for densification at higher temperatures [17]. Densifi cation of SiC-containing ZrB2 powder mixtures initiates at lower temperatures compared with pure ZrB2 as reactions with SiC are deemed to remove the oxide coatings, separat ing ZrB2 particles from mutual contact.  R=3.5mm  40mm  98mm  20mm  Fig. 1. Sharp-shape ZrB2-SiC model used for arc-jet  testing, curvature  radius R = 3.5 mm.  Fig.  2. SEM of  the  polished  surfaces  of  the  ZrB2-SiC ultra  high  temperature ceramics.  X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726  1719  \\x0c', '3.2. Microstructure changes and ablation properties of  the  composite  Characteristics of  the two flat-face models after arc-jet  testing are shown in Table 1. At a heat flux of 1.7 MW/ m2, no measurable weight change was observed, and steady state surface temperature was \\x181650 °C as shown in Fig. 3. However, at a heat flux of 5.4 MW/m2, there was severe  oxidation and ablation after arc-jet  testing and the maxi2300 °C. A total  mum surface  temperatures  exceeded  weight loss of 15.75% was measured and the corresponding  change in thickness was about 3 mm.  Fig. 4 shows  the ZrB2-SiC flat-face model before during (b) and after (c) the ablation test under the heat flux 1.7 MW/m2.  (a),  of  There was  little  change  in  the  visual  appearance of  the  surface and no appreciable  structural  change was detected. The  surface micrograph shows  the  formation of bubbles (Fig. 5), which can be closely related  to the  release of gaseous oxidation by-products  (i.e. CO  and B2O3). The formation of bubbles may imply that diffusion of the formed gaseous phases through the oxide  the  layer  is  slower  than that of O2. to such harsh convectively heating  In fact, during the initial  exposure  conditions,  SiO, B2O3 which tend to evolve outside, most likely coalesced in bub and  other  highly  volatile  boron  sub-oxides,  bles inside the external forming glass [6]. Shear forces asso ciated to the hot stream facilitated the bursting of bubbles  and,  for prolonged exposure,  the smoothing of  the outer most glassy coating. However,  the surface was covered by  a continuous glassy layer. The formation of an external sil ica based glassy layer is very effective in limiting the inward  diffusion of oxygen into the inner bulk, thus enhancing the  Table 1  Summary of arc-jet sample characteristics, conditions and measured surface temperatures  Model  (#)  Dimension  (mm) U30 \\x02 30 U20 \\x02 30  Heat flux (MW/m2)  Weight change  (%)  DThickness  (mm)  Oxide scale (lm)  Ablation time  (s)  Surface temperature (°C)  1  1.7  0.00  0.00  25  600  1640-1660  2  5.4  15.75  2.98  390  600  2150-2330  Fig. 4. Flat-face ZrB2-SiC model before (a), during (b) and after (c) ablation test under a heat flux of 1.7 MW/m2.  0  100  200  300  400  500  600  800  1000  1200  1400  1600  1800  2000  2200  2400  Q=5.4MW/m2  Q=1.7MW/m2  (2)  (1)  T  (  °  C  )  Ablation time (s)  Fig. 3. Temperature curves vs.  time during arc-jet  testing of  the flat-face  ZrB2-SiC models under the different heat fluxes.  Fig. 5. SEM micrograph of the ablated flat-face ZrB2-SiC model under a heat flux of 1.7 MW/m2.  1720  X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726      \\x0c', 'X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726  1721  The macrographs of  ablation test  for 600 s at  the specimen during and after the the heat flux of 5.4 MW/m2 are  shown in Fig. 7. A large amount of  the molten oxidation  product was  blown  away  during  the  ablation  process,  resulting in a high erosion rate and marked change in the  configuration (Fig. 7). The surface of  the ablated sample  was not  smooth and exhibited a mechanical  scour  as a  result of  the airflow. Shear forces associated with the hot  stream facilitated the erosion of  the outermost glass coat ing. It should be noted that the edge of the ablated surface  exhibited a higher erosion rate compared with the central  region. This phenomenon can be explained by the higher  heat flux and shear force at the edge compared with those  at  the center. Moreover,  the high heat flux will result  in a  high temperature at the surface of the sample. Apparently,  the high temperature decreases  the viscosity of glass and  accelerates the chemical reaction,  leading to the increased  oxidation and ablation rates. Fig. 8 is a micrograph of  the ablated surface, which shows the formation of large holes with a diameter of \\x1830 lm. The surface was covered by a limited amount of discontinuous silica glass. Under  this  ablation  condition,  the  SiO2  layer  formed  on  the  Fig. 6. Cross-sectional micrographs of model under a heat flux of 1.7 MW/m2.  the  ablated flat-face ZrB2-SiC  resistance to oxidation. The cross-section of the ZrB2-SiC models after arc-jet testing at a heat flux of 1.7 MW/m2  is shown in Fig. 6. Beneath the external glass layer, there was a subsurface layer about 25 lm thick containing zirco nia. No SiC depletion layer was detected in the present  case.  In fact, during the exposure to such harsh heating,  the models were subjected to mechanical  loading besides  oxidation. Silica alone  cannot develop into a continuous  glassy layer, and will be scoured away by gas flow due to  the existence of the large shear forces during ablation pro cess. The ZrO2 appears to form a porous skeleton that does not enhance the oxidation protection, but it may provide  mechanical integrity and strength to the liquid glass (i.e. sil icate). Moreover,  the ZrO2 skeleton can provide a framethe glass to be retained against the strong gas  work for  flow. The ablation rate is originally controlled by the chem ical reaction of the surface ZrB2 and SiC, and then by the oxygen diffusion through the glass layer and some porous  regions and the evaporation of  the glass. Therefore, most  of the silica glass was retained during the ablation process  and it was distributed in the zirconia grain boundary and  the  surface, which  provides  increased  inhibition  of  the  inward access of oxygen.  Fig. 8. SEM micrograph of the ablated flat-face ZrB2-SiC model under a heat flux of 5.4 MW/m2.  Fig. 7. Flat-face ZrB2-SiC model during (a) and after (b and c) ablation test under a heat flux of 5.4 MW/m2.  \\x0c', '1722  X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726  surface shows a very low viscosity and a significant evapo Underlying this layer was a porous layer, which is a transi ration rate due to the extreme high ablation temperature (>2300 °C), which leads to the intense degradation of  the  tional region between the SiC-depleted layer and external  oxide  layer. The  creation of  such an inner porous  layer  protective effect.  The cross-section of the models after arc-jet the heat flux of 5.4 MW/m2 provided insight  testing at  into micro was most  likely caused by the oxidation of ZrB2 since the formed ZrO2 cannot adhere to the virgin matrix at high temperature. Obviously, cracking  spalling  tend  and  to  structure details of great  interest  (Fig. 9). Three distinct  occur  in this  region due to the weak bonding and CTEs  layers were detected, which is different  from reported lay ered structure in previous studies [18-20]. The outermost layer (350 lm thick) was rather compact except  for a few  large pores. However,  this  layer was not adhered to an  oxide  sub-scale. Compositional analysis by EDS showed  between  mismatch  unaltered ZrB2 matrix. Underneath, a layer partly depleted of SiC (25 lm thick) was formed. The formation of SiC depletion  oxide  scale  and  the  that  the oxide was mainly composed of Zr, O and a little  Table 2  Si. ZrO2 coalesced together by recrystallization under the present condition. Some silica glass was embedded within  the compact ZrO2. The formation of inhibited inward transport of oxygen. Passive oxidation  this layer effectively  protection is provided by the  continuous  layer  that  effectively  prevents  direct  compact ZrO2 exposure of the  ZrB2-SiC composite  to  air  in  this  temperature  region.  Summary  of  arc-jet  sample  characteristics,  conditions  and measured  surface temperatures  Model  (#)  Pressure  (MPa)  Mach  number  Weight  change  Ablation  time (s)  3  4  1.26  1.26  (M)  2.7  2.7  (%)  0.03  8.3  200  100  Surface  temperature (°C)  1445-1451  1443-1450  Fig. 9. Cross-sectional morphologies of the ablated flat-face ZrB2-SiC model under a heat flux of 5.4 MW/m2.  \\x0c', 'X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726  1723  Fig. 10. Photographs of the ZrB2-SiC sharp leading edge model at different testing times.  The ablation conditions for sharp leading edge models  at  the supersonic flow arc-heated wind tunnel are listed in  Zr B2-SiC  C/SiC  Table 2. Fig. 10 shows the macrographs of  (model  3)  sharp leading  edge model  the ZrB2-SiC at different testing  times. The  surface  temperature  increased drastically in a  very  short  time when the  samples were  exposed to the  arc-heated wind tunnel  temperature  rose  dramatically  (Fig. 11). Specifically, 1350 °C,  to  the surface  in mere  3 s,  and then rose gradually, until reached a steady temperature of \\x181450 °C. As a result of the large temperature gradient, a high thermal stress occured. However, no cracks were  observed on the sample. This  fact  suggests  that  the com posite  exhibits  excellent  thermal  shock resistance. There  was a marked change in the color distribution of the model  in the initial 40 s, which then reached a steady state at 60 s,  indicating that the temperature distribution on the surface  of the model achieved a steady state after 60 s of heating. A  similar trend was observed in the C/SiC composite (model  1500  1400  1300  )  C  °  (  T  1200  1100  0  50  100  150  200  Ablation time (s)  Fig. 11. Temperature curves vs.  time during arc-jet  testing of  the sharp  leading edge models.  layer was caused by the active oxidation of SiC, which rep 4) (Fig. 12). But the time to reach steady state of 40 s was  resents a well-known phenomenon in SiC and SiC-contain less than that for model 3. The shorter time to reach steady  ing zirconium diboride-based materials [13,18-20].  state is probably due to its low thermal conductivity com Fig. 12. Photographs of the C/SiC sharp leading edge model at different testing times.      \\x0c', '1724  X.H. Zhang et al. / Composites Science and Technology 68 (2008) 1718-1726  Fig. 13. Photographs of  the sharp leading edge models before ((a) and (c)) and after  ((b) and (d)) ablation testing.  (a) and (b): ZrB2-SiC,  (c) and  (d): C/SiC.  pared with ZrB2-SiC. No configurational and after arc-jet testing was detected for model 3 as shown  change during  In fact, B2O3 has an unusually (450 °C) and a high vapor pressure  low melting  point  (Fig. 14). Therefore,  in Figs. 10 and 13b. This  is quite beneficial  to practical  flight performances. For model 4, C/SiC composite began  to severely ablate at 60 s and structural shape was remark ably  destroyed  as  the  ablation  proceeded. The weight  change was up to 8.3% after arc-jet testing while it was only  0.03% for model 3. Moreover, the structural shape was sig at high temperature, B2O3 vaporizes quickly. The volatilization of B2O3 results in increased oxidation ZrB2 since ZrO2 is not a highly protective oxide. The introduction of SiC remarkably improves the oxidation resis1200 °C due  the ZrB2 material formation of silica glass (Reaction (3) which is more vis above  tance  rapid  of  of  to  the  nificant ablated as shown in Fig. 13d. High oxidation resis cous, having  a higher melting  temperature  and a  lower  tance coupled with configurational stability of  the present  vapor  pressure,  and more  resistant  to  oxygen  diffusion  UHTC when subjected to supersonic flow and heat fluxes  [18-20]. The  typical of a re-entry environment make it a good candidate  for  use  in  aerospace  applications,  especially  for  sharp significant improvement in oxidation resisbased UHTC below 2000 °C has  of ZrB2 achieved by the incorporation of SiC due to the formation  tance  been  shaped structures like wing leading edges and nose caps.  of a silica glass layer with low oxygen permeability, which  provides an efficacious protective oxidation barrier. Apart  3.3. Ablation mechanism  During atmospheric re-entry, space vehicle will undergo  severe aerothermodynamic heating resulting in high tem perature on the outer surfaces of the materials. There were  severe reactions between the reactive gases and the ZrB2- SiC composite. In the present case, the main expected reac tions during the oxidation process are as follows:  O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðlÞ  ZrB2 ðsÞ þ 5 2 B2O3 ðlÞ ! B2O3 ðgÞ  O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ  SiCðsÞ þ 3 2 SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ SiCðsÞ þ 2SiO2 ðlÞ ! 3SiOðgÞ þ COðgÞ SiO2 ðlÞ ! SiO2 ðgÞ  ð1Þ  ð2Þ  ð3Þ  ð4Þ ð5Þ ð6Þ  Fig. 14. Vapor pressure vs.  temperature for the above oxides, calculated  at ambient pressure.  \\x0c', 'from aerothermal  loads, the materials are also subjected to  large aerodynamic loads during hypersonic flight. There fore,  the mechanical  loads  should be taken into account,  as well as oxidation, during practical operations. ZrO2 provides an oxide framework without blown away by the gas  flow leading to the configurational stability of the material.  The combination of ZrO2 and SiO2 is responsible for configurational stability coupled with high oxidation resistance  of  the present UHTC when subjected to simulated atmo spheric  re-entry  conditions. However,  the  efficiency  of  SiC significantly decreased at higher 2200 °C,  temperatures,  espe cially  above  because  of  the  active  oxidation  (Reactions  (4) and (5) and rapid evaporation (Reaction  (6)) due to high vapor pressure (Fig. 14). Moreover,  the  existence of  the  large  amount of  silica will  significantly  lower  the melting  point  of ZrO2-SiO2, which the formed oxide scale resulting in spalla causes  a  strength loss of  tion or peeling under aerodynamic loads.  Apparently,  the  present UHTC cannot maintain  the  passivating protective capability at extremely high temper(i.e. >2300 °C)  atures  since the formed oxide scale is not  stable under the aerothermodynamic loads. The tempera ture limit  for SiC reinforced ZrB2 based UHTC is mostly dependent on the softening and degradation of ZrO2 based oxide scale, which is also controlled by SiC content since  the melting point of  the ZrO2 based oxide scale is closely connected with SiO2 content. When the shear forces exceed the strength of the formed oxide framework, the oxides  peel  off  or  dislodge  from the  specimens  resulting  in  increased  oxidation  and  degradation  of  configuration,  which in turn affect the performance and safety of the vehi cles. The softening and degradation of oxide scale resulted  from increased temperatures is a major issue of the UHTC  for use in extreme environments like those associated with  hypersonic flight, atmospheric re-entry and rocket propul sion. In this respect,  the content of  the low melting point  oxides can be reduced to improve the upper limit of the ser vice  temperatures. High SiC content  is beneficial  for  the  performance of  the UHTC at  low temperatures, whereas  low SiC content  is favorable for the use of  the UHTC at  high temperatures. Therefore,  the  compositions of  rein forced phases (i.e. SiC) should be optimized according to  the  practical  service  environment, with  damage within  acceptable limits.  The coherence between oxide scale and unaltered mate rial  is another critical  issue. As can be seen from Fig. 9,  many defects like pores existed in the inner material which  may  lead to the  rupture of  the whole oxide  scale. The  extent of the damage under these conditions is to be consid ered unacceptable for  the foreseen demands of  reliability  and re-usability at extremely high temperatures (i.e. above 2300 °C).  4. Conclusions  An ultra high temperature dense ZrB2-SiC ceramic was produced by hot-pressing. ZrB2-SiC models were exposed  to ground simulated atmospheric re-entry conditions using arc-jet testing with heat fluxes of 1.7 MW/m2 and 5.4 MW/ m2, respectively. For temperatures in the order of 1600- 1700 °C, the material was able to endure the heating condi tions,  thanks  to the formation of an external multiphase  oxide scale. However, for temperatures in the order of 2300 °C, evident oxidation and ablation occured and the  material was unable to offer a valuable  resistance  to the  applied aerothermal load. ZrB2-SiC ultra high temperature ceramics also exhibited an excellent thermal-oxidative and  configurational stability under supersonic conditions com pared with the C/SiC material, which suggests  they are  potential candidates for leading edges. Results indicate that  ZrB2-SiC can maintain the high oxidation resistance coupled with configurational stability at temperatures lower  than that point which results  in significant  softening and  degradation of  the oxide scale, and that point will be the  temperature limit  for UHTC.  Acknowledgements  This work was supported by the National Natural Sci ence Foundation of China (90505015 and 50602010),  the  Research Fund for the Doctoral Program of Higher Edu cation (20060213031) and the Program for New Century  Excellent Talents in University.  References  [1] Kuriakose AK, Margrave JL. The oxidation kinetics of  zirconium  diboride and zirconium carbide at high temperatures. J Electrochem  Soc 1964;111(7):827-31.  [2] Tripp WC, Davis HH, Graham HC. Effect of a SiC addition on the  oxidation of ZrB2. Am Ceram Soc Bull 1973;52(8):612-6.  [3] Hinze  JW, Tripp WC, Graham HC. High-temperature oxidation  behavior of a HfB2 plus 20 v/o SiC composite. 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[20] Fahrenholtz WG. Thermodynamic analysis of ZrB2-SiC oxidation:  Processing, properties  and arc  jet oxidation of hafnium diboride/  formation of a SiC-depletion region. J Am Ceram Soc 2007;90(1):  silicon  carbide  ultra  high  temperature  ceramic.  J Mater  Sci  143-8.  2004;39(19):5925-37.  [16] Han JC, Hu P, Zhang XH, Meng SH, Han WB. Oxidation resistant  ZrB2 -SiC Composites at 2200 C. Compos Sci Technol 2008;68(3-4):  799-806.  \\x0c']"
},{
  "_id": 8,
  "PDF": "Ablation behaviourofaTaC coating on SiC coated C C composites at different temperatures.pdf",
  "Text": "['CERAMICS  INTERNATIONAL  Available online at www.sciencedirect.com  Ceramics International 39 (2013) 359-365  Ablation behaviour of a TaC coating on SiC coated C/C composites  at different  temperatures  Yong-jie Wang, He-jun Lin, Qian-gang Fu, Heng Wu, Lei Liu, Can Sun  State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, PR China  Received 7 May 2012; received in revised form 6 June 2012; accepted 11 June 2012  Available online 17 June 2012  Abstract  To improve the ablation resistance of carbon/carbon (C/C) composites, a TaC coating was prepared by supersonic plasma spraying  on SiC coated C/C composites. The microstructure and morphology of the coatings were characterised by Scanning Electron Microscopy and X-ray diffraction. The ablation properties were studied at different temperatures under oxyacetylene torch. At 2100 1C, the oxides were blown away and resulted in high ablation rates: 1.2 \\x02 10 \\x00 2 mm/s and 3.9 \\x02 10 \\x00 3 g/s. However, most oxides can remain in ablation centre and serve as a coating at low temperature (1900 and 1800 1C). Therefore, the TaC/SiC coated samples exhibited zero  linear ablation rate and lower mass ablation rate.  & 2012 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  Keywords: B. Composites; C. Corrosion; E. Thermal applications  1.  Introduction  Carbon/carbon composites (C/C) are appropriate struc tural materials in ultra-high-temperature environments due  to their outstanding  thermal properties  [1-3]. However,  without protection these composites will be damaged by ablation at high temperature above 2000 1C [4]. In order to  improve  the  ablation resistance of C/C composites  and  ensure the composites have a longer  serve life,  refractory  carbides have been introduced into C/C composties  [5].  TaC has  always  been  concerned  because  it  has  some  remarkable properties such as high hardness, high melting point (above 3880 1C) and good thermal conductivity [6,7].  Up to  now,  two ways  are  commonly used to improve  ablation resistance of C/C composites by TaC. One way is  to add TaC into the matrix of  these  composites. Xiang  et  al.  [8]  prepared Tantalum carbide materials  in C/C  composites by  liquid precursor  conversion method. The  TaC grain size enlarged with temperature elevating. Wang  et al.  [9]  introduced TaC powders  into C/SiC composites  using slurry infiltration method. The glass state Ta2O5 can  improve anti-ablation performance at ultra-high tempera ture. Chen et al.  [10] fabricated C/C composites with SiC-  TaC  inter  layer  through  isothermal  chemical  vapour  infiltration. But,  the  tantalum compounds were not able  to seal off the material surface during ablation. A uniform,  adherent  and  crack-free TaC coating was  obtained  on  carbon fibres using a molten salt method by Dong et al.  [11],  and thermo-gravimetric  analysis  indicated that  the  oxidation  resistance  of  carbon  fibre  can  be  improved  remarkably  with  a  high-quality  TaC  layer.  In  these  methods, TaC served  as  additives  in matrix  to  abate  ablation of C/C composites. However, TaC phase cannot  be homogeneously introduced, which has a negative effect  on the mechanical properties of  the composites. Further more,  it is limited to improve the ablation resistance of C/  C composites and the modified composites still show high  ablation rates,  through these methods.  The other way is to prepare TaC coating on the surface  of C/C composites. TaC coating was  fabricated on C/C  surface using ethylate tantalum as precursor by He et al.  [12], the coating had single treatment above 1400 1C. But,  phase  composition  after  the coating seemed loosen  and large amounts of defects (cracks, pinholes) existed in  the  coating, which are disadvantageous  for  the ablation  www.elsevier.com/locate/ceramint  0272-8842/$36.00 & 2012 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  http://dx.doi.org/10.1016/j.ceramint.2012.06.034  nCorresponding author. Tel.: þ 86 29 88495004;  fax: þ 86 29 88492642.  E-mail address:  lihejun@nwpu.edu.cn (H.-j. Li).  \\x0c', 'resistance. Li et al. [13], Chen et al. [14] have prepared TaC  coating by means of  chemical vapour deposition (CVD)  and developed several TaC coating systems on C/C sur face. These  coatings have dense  surface, and performed  better ablation properites. But,  the CVD is an expensive  method. Moreover,  the  controlling of TaCl5  sublimation  process needs high requirements of  the  equipment. Low  temperature phase of Ta2C is always accompanied in the  TaC coating, which has  adverse  effects on the  ablation  resistance.  In this work, supersonic plasma spraying was employed  to prepare a TaC coating on SiC coated C/C composites. It  is a simple and low-cost way. Compared with traditional  plasma spraying, the plasma temperature is 10000 1C and it has a jet velocity up to 600 m/s  above  [15,16].  Dense  coating with better bond can be obtained by this  method. The phase  compositions and microstructures of  the  as-prepared coating were discussed. Meanwhile,  the  coating ablation test was carried out at different  tempera tures,  in order to study the oxides state during ablation.  2. Experimental  2.1. Preparation of TaC/SiC coated C/C composites  The substrates (F 30 mm \\x02 10 mm) were cut  from bulk  2-D C/C composites.  The  composites were  fabricated  through isothermal chemical vapour infiltration with a final density of 1.78 g/cm3. Then, the specimens were hand abraded using 300 grit SiC paper, water and dried at 200 1C for 2 h.  cleaned with distilled  The as-received samples were embedded in mixed powder  in a graphite crucible, of which the composition is 65 wt% Si,  20 wt% graphite and 15 wt% Al2O3. The crucible was placed into furnace and held at 2000-2200 1C for 2 h to obtain a SiC  coating. Details have been reported in Ref.  [17].  Then TaC coating was prepared on SiC coated compo sites through supersonic plasma spraying. TaC powder (purity 4 99.9%) was provided by ZhuZhou GuangYuan  Cemented Material Co., Ltd. The size of TaC grains was 1-1.5 mm. The argon was employed as the primary gas and  the carrier gas. The hydrogen was used as  the secondary  gas. Details  of  the  spraying  parameters  are  listed  in  Table 1. The details of  the spraying have been described  in Ref.  [15].  2.2. Ablation tests and microstructure analysis  Ablation  behaviour  was  tested  under  oxyacetylene  torch. The TaC/SiC-coated samples were placed vertically  to the flame. The inner diameter of  the nozzle tip of  the  ablation gun was  2.5 mm. The  surface  temperature was  detected through optical pyrometer.  In our work,  three  different ablation temperatures were employed in the ablation (2100, 1900 and 1800 1C). The temperatures were (1872 1C),  chosen around the melting point of Ta2O5  in  order to analyse the fusant state after ablation. The linear  and mass ablation rates of  the samples could be obtained  according to Eqs.  (1) and (2) below.  Rl ¼ Dd =t  ð1Þ  Rm ¼ Dm=t  ð2Þ  Rl  is linear ablation rate; Dd is the change of  the sample’s  thickness at centre region before and after ablation; Rm is mass ablation rate; Dm is sample’s mass change before and  after ablation;  t  is ablation time.  The ablation morphology, microstructure and elemental  composition of TaC/SiC coated composites were examined  by  scanning  electron microscope  (FE-SEM SUPRA-550).  The phases of the coating were characterised by Rigaku D/max-3C X-ray diffraction (XRD, 40 kV, 35 mA, Cu Ka).  3. Results and discussion  3.1. Microstructures of  the coating  From Fig. 1,  it can be seen that the coating is composed  of TaC and Ta2O5. During  spraying, the plasma 10,000 1C,  could  offer  high  temperature  above  and  the TaC  powder would be quickly melted under the plasma torch.  However,  the melted TaC can  be  oxidised  into Ta2O5  according Eqs.  (3) and (4). Therefore, Ta2O5 phases are  detected in the XRD pattern. 2TaCðsÞ þ 7=2O2 ðgÞ-Ta2O5 ðsÞ þ 2COðgÞ  ð3Þ  2TaCðsÞ þ 9=2O2 ðgÞ-Ta2O5 ðsÞ þ 2CO2 ðgÞ  ð4Þ  Fig. 2 is the surface and cross section BSE images of the  TaC/SiC coating. The coating is covered by lots of white  mottles, and no cracks are found on the surface (Fig. 2a).  From the magnified image (Fig. 2b), the white mottles can  be  recognised as  fusing state,  from which the  formation  mechanism of the coating can be inferred. During ablation,  the TaC was  fused in a  short  time. The melted grains  crashed and spread out on the surface of substrate, formed  the TaC layer. This process was  repeated and finally  a  thick coating was obtained. In spraying, gases (such as air,  CO or CO2) may be mingled into the powder,  so small  pinholes can be observed at the surface. Owing to the high  speed of  the plasma flame,  the oxidation of TaC and the  mingled gases were much less  than that of  the normal  speed spraying. That  is the reason why supersonic plasma  spraying can get a dense  coating (Fig. 2c). The SiC and  Table 1  The spraying parameters for the TaC coating.  Spraying voltage (V)  360-400  Spraying current  (A)  140-170  Spraying distance (mm)  90  Primary gas (Ar) flow rate (L/min)  80-100  Secondary gas (H2) flow rate (L/min)  30-40  Carrier gas (Ar) flow rate (L/min)  6-10  Powder feeding rate (g/min)  about 35  Nozzle diameter (mm)  4.5  Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  360  \\x0c', 'TaC coatings can be distinguished easily, with thickness of about 50 and 250 mm respectively.  3.2. The ablation behaviour of  the TaC/SiC coating  From Fig. 3a,  it can be found that the surface tempera ture rises quickly in the first 30 s for the 2# and 3# curves.  After 30 s the temperature increases very slowly and finally  stays almost at a fixed value, which is  close  to our  test  temperature.  It  is because  that  a heat balance may be  achieved between the heating  (from oxyacetylene flame)  and  cooling  processes  (heat  emission  and  transmission  from water cycling system). Thus, the surface temperature  will  stop to rise. Meanwhile,  the  evaporation of Ta2O5  consumed a large amount of heat, which could lower the surface temperature. As to the three temperatures, 2100 1C and 1900 1C are above  the melting point of Ta2O5, the than that of 1800 1C. So, flat  evaporation is more fierce  steps can be found, which did not appear temperature curve at 1800 1C.  in the surface  From the macro  images  (Fig.  3b-e)  of  the  ablated  samples, great differences happen to the samples’ surface. 2100 1C,  At  deep  pits  exist  in  the  central  region;  the  coating has been broken through after ablation for 60 s; image. At 1900 1C,  the black matrix can be  seen in the  fusant appeared in border  region, but the coating is 1800 1C, most  still  integrated  in  centre  region. At  fusant  assembled in centre  region and no fusant was blown to  the border region. 2100 1C,  At  the  temperature  is much higher  than the  melting point of Ta2O5. The  fusant had higher  liquidity  and lower viscidity, The coherence between the fusant and  coating became very poor. Most  fusant was blown away  by the flame and little remained on the surface. The fusant  would not behave positive  ability  to the  ablation resis tance. Therefore, the composites were attacked severely by  the flame, deep pits  can be  found in centre  region. The  coated  composites  have  bad  ablation resistance (1.2 \\x02 10 \\x00 2 mm/s  and  showed 3.9 \\x02 10  high ablation \\x00 3 g/s, Fig. 4) at  rates  and  this  temperature. However,  the  fusant had a higher viscidity and lower and 1800 1C. They can adhere  liquidity at 1900  to the  inner  coating and  Fig. 1. XRD patterns of as-received TaC/SiC coating.  50µm  100µm  5µm  Fig. 2. Surface and cross section SEM images of TaC/SiC coating.  Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  361  \\x0c', 'provide efficient protection. Both the linear ablation rates ( \\x00 3.58 \\x02 10 \\x00 3 mm/s and \\x00 2.6 \\x02 10 \\x00 3 mm/s),  are minus  that  is,  the  coating  behaved zero 1.7 \\x02 10 \\x00 3 g/s  ablation. The mass 0.48 \\x02 10 \\x00 3 g/s,  ablation  rates  are  and  respectively, which are also much lower than that at 2100 1C. At 1800 1C, almost no fusant was blown away,  therefore TaC coating behaved lowest mass ablation rate.  Fig. 5a and b are the images after 60 s ablation at 2100 1C. The fibres were burned into cone shape. Some where  the fibres were  cut off by flame  shear  force, only  holes were left. No oxides can be found on surface. It can  be  inferred that  the  coating has  failed completely. The  irregular shaped pits on the matrix and fibres verified that  the composites suffered severe ablation under oxyacetylene  torch.  In  order  to  research  the  oxides  state  change,  the  morphology in centre region for 20 s ablation is presented  in Fig. 5c and d. The coating has been broken through, the  matrix cannot be protected. The fusant presented not as a  coating but spheres, which covered on the fibres. The fibre  appeared as rough surface. Additionally,  some hemisphe rical  pits  can  be  found  on  fibre  surface.  From the  magnified photo (Fig.  5d),  tiny  residual Ta2O5 or SiO2  spheres can be found at  the pit bottom. It  is deduced that  the  oxides may  accelerate  the  corrosion  to  the  fibres.  During ablation,  the evaporation and melt of Ta2O5 and  SiO2 can consume a lot of heat, it is positive to the ablation  protection. However, unlike the positive effect of oxides on  the ablation resistance in matrix modified C/C composites  [18],  the  oxides  cannot  be  provided  continuously.  Its  positive  effect  is  very  slight. The  oxides were mostly  inclined  to  accelerate  the  corrosion  to  the  fibres.  It  is  because the oxides reacted with carbon matrix according to  Eqs.  (5)-(8). After  that,  the TaC and SiC continued to  react with oxygen and generate oxides again (Eqs. (3), (4),  (9)  and (10)).  In this process,  the  carbon was  the only  consumed substance. The  residual oxides  just  served as  catalyst and did not consume according to these reactions.  The acceleration would not stop until the residual oxides were blown out. The oxides were spherical at 2100 1C, they  corroded the fibres  so that  the pits presented as hemi sphere-shaped. When the big spheres were blown away,  tiny ones remained in the pits bottom. Ta2O5 ðl Þ þ 7CðsÞ-2TaCðsÞ þ 5COðgÞ  ð5Þ  Ta2O5 ðl Þ þ 9=2CðsÞ-2TaCðsÞ þ 5=2CO2 ðgÞ  ð6Þ  SiO2 ðl Þ þ 3CðsÞ-SiCðsÞ þ 2COðgÞ  ð7Þ  SiO2 ðl Þ þ CðsÞ-SiCðsÞ þ CO2 ðgÞ  ð8Þ  SiCðsÞ þ 2O2 ðgÞ-SiO2 ðl Þ þ CO2 ðgÞ  ð9Þ  SiCðsÞ þ 3=2O2 ðgÞ-SiO2 ðl Þ þ COðgÞ  ð10Þ  Fig.  6  shows  the SEM images of centre 1900 1C. Much fusant  region after  ablation for  60 s  at  appeared on  surface. However, Ta2O5 can still be blown away but  just  to the border region at this temperature. From Fig. 6a, the  ablation mechanism of the coating can be inferred. During  ablation,  the bulge place on surface is more close to the  nozzle  tip  and  has  higher  temperature  than  the  other  places. So, Ta2O5 melted firstly there. Because of mechan ical denudation of  the flame, partial Ta2O5 was blown  away.  Then,  an  ablation  pit  formed  there,  which  is  presented in the centre of Fig. 6a. However,  the pits edge  became into bulge place, and underwent the same process,  as  ablation  going.  The  circulation  repeated  until  the  Fig. 3. The surface temperature and the macro images of the coated samples at different temperature. (a) (1#:1800 1C, 2#:1900 1C, 3#:2100 1C), (b) :1# (c) :2#  (d)  :3# (e) (ablation region sketch.  (C): centre region,  (B):  transition region, (A): border region).  Fig. 4. The ablation properties of coated samples at different temperatures.  Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  362  \\x0c', 'Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  363  10µm  5µm  5µm  1µm  Fig. 5. SEM image of ablation centre region at 2100 1C (a and b) 60 s,  (c and d) 20 s.  pits   cracks  50µm  5µm  Fig. 6. SEM images of centre region at 1900 1C.  coating was broken through. According to the 1900 1C,  exhibited  good  the  coating  rates  at  ablation  ablation  resistance.  It  can  be  concluded  that  the mechanical  denudation did damages  to the coating, but  it could not  ruin the integrality of  the coating and affect  the ablation  resistance. From Fig. 6b,  it  can be  seen that  the  fusant  existed as dendrites after ablation.  It  is because  that  the  fusant  is  composed  of monoclinic Ta2O5  (Fig.  7). The  phase has a high viscidity and is apt  to form branch-like  crystalline, when cooled down from high temperature. The  dendrites have  large  size. Cracks  can  be  found  on the  surface, which may provide channels for oxygen and result  in oxidation of  inner coating or composites.  However,  the  surface  is much  different  at  1800 1C  (Fig. 8). The surface is flat, some micro holes can be found  Fig. 7. The XRD pattern of centre region after ablation.  and  the  fusant  distributes  uniformly  in  centre  region  found on the surface (Fig. 8b). The melting Ta2O5 formed  (Fig.  smaller  8a). The size of branch-like crystalline than that at 1900 1C, and no obvious  is much  a dense coating on the surface. It  is advantageous for the  cracks are  ablation resistance. The melting  coating  can serve  as  a  \\x0c', 'thermal barrier  layer and bring down the inner  tempera ture of  the coating. Meanwhile,  the melting Ta2O5 had a  low oxygen-diffusing coefficient and prevented inner coat ing from oxidation. Therefore, the coating exhibited better  ablation resistance at  this temperature.  It  can  be  concluded  that  the  state  of  oxides  during  ablation is very important  for ablation resistance proper ties of  the TaC coating on SiC coated C/C composites.  When  the  temperature  is  just  below the  oxide melting  point,  the oxide  fusant has a proper viscidity to form a  glassy  coating  on  the  surface. The  glassy  coating  can  provide  a  thermal  and oxygen diffusing barrier  for  the  inner coating. The mechanical denudation and the oxida tion mechanism can be held furthest. If the temperature is  much higher  than oxide melting point,  the viscidity is  so  low and the melting oxides  can be blown away through  mechanical  denudation.  The  coating  will  be  broken  through rapidly.  4. Conclusions  A TaC/SiC coating was  prepared  through  diffusing  reaction and supersonic plasma spraying method. Surface  temperature  stayed  at  a  fixed  value  after  30 s  ablation  under oxyacetylene torch, owing to the oxides evaporation.  The TaC/SiC coating showed better ablation resistance at 1800 1C. The  1900  and  oxides  can  remain  in  ablation  centre because of their high viscidity. The oxides served as  coating during ablation and provided efficient protection for composites. At 2100 1C,  the viscidity of oxides was so  low that  the melts were blown away by the flame. More over,  the residual oxides accelerated the corrosion to C/C  composites after the coating was broken through.  Acknowledgements  This work has been supported by the National Natural  Science Foundation of China under Grant Nos. 51072166  and 50902111,  the Program for New Century Excellent  Talents  in University,  the Research Fund of State Key  Laboratory of Solidification Processing (NWPU), China  (Grant No.  25-TZ-2009) and ‘‘111’’ Project  (Grant No.  B08040).  References  [1]  J.D. Buckley, D.D. Edie, Carbon—Carbon Materials and Compo sites, Noyes Publication, 1993 (pp. 267-279).  [2] C.R. 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Huang, Structural character istics and formation mechanisms of  crack-free multilayer TaC/SiC  coatings on carbon-carbon composites, Transactions of Nonferrous  Metals Society of China 18 (2008) 255-261.  [14] Z.K. Chen, X. Xiong, G.D. Li, W. Sun, Y. Long, Texture structure  and ablation behaviour of TaC coating on carbon/carbon compo sites, Applied Surface Science 257 (2010) 656-661.  [15] H. Wu, H.J.  Li, Q.G.  Fu, D.J. Yao, Y.J. Wang, Chao Ma,  Microstructures and ablation resistance of ZrC coating for SiC-coated  50µm  5µm  Fig. 8. SEM images of centre region at 1800 1C.  Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  364  \\x0c', 'Yong-jie Wang et al. / Ceramics International 39 (2013) 359-365  365  carbon/carbon composites prepared by  supersonic plasma  spraying,  [17] K.Z. Li, D.S. Hou, H.J. Li, Q.G. Fu, G.S. 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},{
  "_id": 9,
  "PDF": "Ablation characteristics of rocket nozzle using HfC-SiC refractory ceramic composite.pdf",
  "Text": "['Acta Astronautica 173 (2020) 31-44  Contents lists available at ScienceDirect  Acta Astronautica  jou rna l homepage : www .e lsev ie r .com / loca te /ac taas t ro  Ablation characteristics of rocket nozzle using HfC-SiC refractory ceramic composite  T  Kyu-Seop Kima, Sea-Hoon Leeb, Van Quyet Nguyenb, Yongtae Yuna, Sejin Kwona,∗  a Korea Advanced Institute of Science and Technology, Department of Aerospace Engineering, KAIST, 291 Daehak-ro, 34141 Daejeon, Republic of Korea b Korea Institute of Material Science, 797 Changwon-daero, 51508 Changwon, Seongsan-gu, Republic of Korea  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ablation Ultra High-Temperature Ceramics (UHTCs) Hybrid rocket Oxidation Graphite Hydrogen peroxide  Modern hybrid rockets are susceptible to the highly ablative environment in which they operate, as introducing an adequate cooling system has so far been a challenge. Ablative material has been widely adopted for use in hybrid rockets, which can only survive for the limited operation time. The use of ablative materials such as pyrolytic composite or graphite is advantageous for use in a simple system in an uncooled state. However, ablation causes severe enlarging in the nozzle throat area and a significant drop in the pressure during operation. Recently, Ultra High-Temperature Ceramics (UHTCs) have been noted as superior in terms of ablation and oxidation resistance. The most promising HfC-SiC composite ceramic has demonstrated an enhanced performance for thermal and mechanical properties. In this study, an HfC-SiC composite was embedded in the graphite casing as a nozzle insert throat to analyze the feasibility of its application in rockets. An experimental analysis of the ablation in the HfC-SiC was carried out with a 250 N scale hybrid thruster using High Test Peroxide (HTP) and High-Density Polyethylene. The hot-fire condition was set to above 30 bar for 25 s combustion, with the purpose of producing significant erosion on the nozzle materials. The graphite nozzle of the same shape was also tested under the same experimental conditions for comparison of the erosion. The hot-fire test with the HfC-SiC insert resulted in a stable rocket performance, with improvements to the chamber pressure and the thrust, whereas the combustion performance varied undesirably in the test using the graphite nozzle due to ablation. The rate of ablation in the throat was significantly reduced to 15.81 μm/s using HfC-SiC, which is 46.5% of the erosion rate found in the graphite. The adoption of the HfC-SiC composite effectively inhibited the ablation on the throat. However, substantial erosion occurred on the interfaces due to the different ablation resistances. The feasibility of adopting hafnium-based carbide on the rocket was also evaluated based on a material analysis of the surface oxidation.  1.  Introduction  Hybrid rocket systems are attractive for use as propulsion systems because of the safe handling, easy use, and simple schematics. The hybrid rocket has been actively researched for application as the nextgeneration launch vehicle. A significant amount of research into enhancing the performance of the hybrid rocket has been carried out. One of the state-of-the-art technologies in the hybrid propulsion system is the prevention or reduction of ablation on the nozzle throat. As the boundary layer thickness tends to be the minimum at the vicinity of the nozzle throat, maximum heat flux and shear stress are inherently exerted on the throat [1]. Therefore, this extreme condition generally leads to the destruction of the nozzle surface, especially the throat. The implementation of cooling in hybrid thrusters to reduce erosion rate is  complicated, unlike in liquid propulsion where cooling methods using propellants for both the chamber and nozzle are available. For this reason, ablative materials such as graphite and pyrolytic composite has been adopted for use in the general hybrid motor. The ablative material is simple to use and can withstand extremely high-temperatures in an uncooled state. However, these materials suffer significantly from oxidation and thermal decomposition, leading to the occurrence of ablation during operation. Ablation is known as the main factor that induces a pressure drop in the combustion chamber and poor combustion [2], and is a result of the combined effects of thermochemical and thermo-mechanical erosion. Many numerical studies have estimated the thermochemical ablation for ablative materials [3-6]. L. Kamps et al. [7] investigated throat-erosion history by the reconstruction method and reported trends in sync with the numerical estimation  ∗ Corresponding author. E-mail addresses: kswind@kaist.ac.kr (K.-S. Kim), seahoon1@kims.re.kr (S.-H. Lee), quyet_win@kims.re.kr (V.Q. Nguyen), straw_hat@kaist.ac.kr (Y. Yun), trumpet@kaist.ac.kr (S. Kwon).  https://doi.org/10.1016/j.actaastro.2020.03.050 Received 1 November 2019; Received in revised form 20 February 2020; Accepted 29 March 2020 Available online 07 April 2020 0094-5765/ © 2020 IAA. Published by Elsevier Ltd. All rights reserved.  \\x0c', \"K.-S. Kim, et al.  Nomenclature  throat area empirical constant characteristic velocity fuel port diameter thrust thrust of initial stage oxidizer mass flux Gibbs' free energy length of fuel grain total propellant flow rate oxidizer mass flow rate fuel mass flow rate choked mass flow rate mass flux exponent oxidizer-to-fuel mass ratio  Acta Astronautica 173 (2020) 31-44  chamber pressure reversely-estimated chamber pressure using minimum insert radius reversely-estimated chamber pressure using average insert radius universal gas constant regression rate temperature inner nozzle radius before combustion test inner nozzle radius after combustion test ablated radius ablation rate burn time pressure rising time specific heat ratio fuel density characteristic velocity efficiency  [6]. However, thermomechanical erosion also plays a significant role in rocket combustion [8,9]. Thermomechanical erosion is extremely difficult to predict numerically and can only be analyzed via experiment. It is therefore difficult to predict variations in the rocket performance, such as the drop in pressure that occurs in the chamber, when the ablation on the nozzle is severe. The proper prediction of rocket performance requires numerous sets of experiments and a significant period of time. If the nozzle material can withstand the high pressure and temperatures in an uncooled state and with negligible ablation, it is possible that the hybrid motor could overcome the limitations in operation that are brought about by the destruction of the material. This enables stable combustion without a severe degradation in performance. Therefore, it is essential to adopt pertinent materials for use in the hybrid rocket system, and alternative ablative materials need to be refractory and resistant to ablation in the environment of rocket combustion in an uncooled state. A few studies were carried out in the 1960s concerning this matter, and candidate ceramics for use in rocket nozzles have been found [2,10]. Several monolithic ceramics have been tested in solid rocket motors, with extensive research conducted on pyrolytic refrasil [11,12]. However, this refractory ceramic was found to be affected by thermal shock. Pyrolytic refrasils such as silica phenolic are only operable for a limited time in the rocket environment because of the substantial thermal decomposition and surface melting that occurs [13]. Recently, Ultra-High Temperature Ceramics (UHTCs) have been noted to be highly refractory along with superior oxidation resistance. The promising candidates of the UHTCs have been suggested as advanced alternatives for extreme environment applications [14-16]. Some of these candidates showed superior oxidation and ablation resistance characteristics at elevated temperature environment [17-22] and in propulsion environment [23,24]. However, there are only a few studies on testing the UHTC as a rocket nozzle in a hybrid thruster. A few materials (such as TaC-based, SiC-based, and fiber-reinforced ZrB2SiC) have been tested, and these have shown enhanced ablation resistance at the throat [25-27]. Ceramic compounds of refractory transition metals, especially carbide, have incomparably superior refractory characteristics. Among the carbide group, hafnium carbide ceramics are the most promising. Hafnium carbide has a melting point of over 3273 and highly stabilized chemical properties in extreme temperatures. Unfortunately, monolithic UHTCs, including refractory carbide, suffers from thermal shock because of its brittleness. Possible methods for improving the thermal shock resistance have been suggested, some of which include reinforcing the ceramics with fiber and addition of silicon carbide. D.  Sciti et al. [17] and S. Mungiguerra et al. [23] tested fiber-reinforced ceramic materials in rocket applications to resolve the low fractural toughness of the UHTC. Furthermore, the addition of silicon carbide to the Hf-based ceramic can not only enhance its thermal shock resistance, but is also resistant to oxidation as it forms SiO2, which hampers the oxygen diffusion with HfO2 [15,18]. Thus, the HfC based composite with the addition of the SiC has been considered as the suitable target material for application in rockets. The HfC-SiC composite is expected to have a self-generating oxide layer when oxidized. This oxide layer would efficiently hamper the diffusion of the active oxygen and protect the interior material by forming a dense layer, which buffers the ablation between the combustion gases and the interior material [19]. The present authors also tested the oxidation and ablation behavior of HfC-SiC nano-composites with 20 wt % HfSi2-C additives and reported the excellent properties of the materials above 2773 K using an oxy-acetylene torch for an extended period over a range of 5-30 min [28]. The oxidized composite was found to form a dense oxide layer on the surface and demonstrate a linear ablation rate of approximately 0.14 μm/s, which implied an enhanced oxidation resistance compared to other UHTC candidates. The temperature of the common combustion gases at the rocket is expected to reach over 2500 K; comparable to the temperature of the oxyacetylene torch test [29]. However, the surface erosion of the HfC-SiC composite is considered to be more significant under a high-pressure rocket environment even if the thermochemical effects on the ablation from the composition on the different oxidizing compound are excluded. The presence of high-pressure gases exerts harsh shear stresses on the nozzle surface, meaning that thermomechanical ablation may play a critical role in the erosion of the HfCSiC. Thus, this study was carried out to achieve two purposes: validation of the feasibility of adopting hafnium carbide-based ceramics into the rocket nozzle and investigating the ablation characteristics of the HfCSiC composite under a rocket combustion test. A comprehensive analysis was conducted based on the performance of the rocket, a quantitative measurement of the ablation using LASER 3D scanning, and the surface morphologies. A hybrid thruster, using High-Test Peroxide (HTP) and High-Density Polyethylene (HDPE), was used for producing substantial erosion of selected target materials, which were the HfC-SiC insert and the graphite. As previously discussed, a severe throat ablation was expected due to the extreme heat flux and shear stress. The authors intended to insert the HfC-SiC composite into the throat to enable it to withstand harsh conditions. Since it is known that ablation becomes more severe as the pressure in the chamber pressure increases, a static hot-fire test condition was set to a reasonably high pressure for an adequate time period for the production of surface erosion.  32  KAtaCDpFFiGoxGLmtmoxmfmchokednO/FPcPrevmin,Prevavg,RrTaftbfrbprfc\\x0c\", \"K.-S. Kim, et al.  2. Experimental methodology  2.1. Hybrid thruster design  To test the ablation characteristics of the HfC-SiC composite, a 250 N scale hybrid thruster was designed, using 90 wt % hydrogen peroxide (HTP) and High-Density Polyethylene (HDPE). The experimental setup is presented in Fig. 1. The HTP oxidizer was pressurized with gaseous nitrogen at a pressurization facility, and an accumulator was mounted for suppressing the pressure surge in the early transient period. The mass flow rate of the oxidizer was measured using the differential pressure of the mass flow meter (MFM) which was carefully calibrated before the combustion test. The hybrid thruster was constituted of a shower-head type injector, a catalyst bed containing the catalyst used for decomposing the HTP, a combustion chamber with fuel grain, and a nozzle. The HTP was injected into the catalyst bed through the shower-head. Ignition of the fuel grain occurred automatically without external ignition source, as the highly concentrated HTP decomposed and generated the heat required for the combustion. The temperature of the decomposed HTP gases increased to approximately 1022 K while passing through the catalyst chamber, leading to the auto-ignition of the HDPE fuel grain, which has a natural ignition temperature of approximately 570-720 K. The MnO2/La/Al2O3 catalyst was prepared using the method proposed by Lee et al. [30] for the efficient decomposition of the HTP even with 25 s of combustion. The chamber pressure and combustion time were set to 30 bar and 25 s, respectively, to induce significant ablation. A comparison of the characteristics of the erosion with the widelyused ablative graphite was necessary for validating the superior ablation resistance of HfC-SiC. Therefore, the selected target materials were graphite and HfC-SiC. The nozzle structures were then prepared using the following procedures.  2.2. Preparation of nozzle specimen  A commercial HfC (D50: 1.1 μm; Japan New Metal, Osaka, Japan), HfSi2 (particle size: ~30 μm, purity: > 99%; Alfa Aesar, MA, USA) and  Acta Astronautica 173 (2020) 31-44  carbon black powers (surface area: 60-80 m2/g, purity: 99.5%; Alfa Aesar, MA, USA) were used as starting materials. Stoichiometric amounts of carbon black and HfSi2 and were used as the sintering additives of HfC. The additives finally reacted to form HfC and SiC in the final composites according to the following Eq. (1): [31].  (1)  where ΔG is the Gibbs' free energy in J/mol, and T is the temperature in K. The mixing ratio of HfSi2-C additives to HfC was fixed as 5 wt %. The starting materials were mixed using a high energy ball milling machine (HEBM, SpexD8000, Spex CertiPrep, Metuchen, NJ) for 2 h under dry condition with WC balls and jars. The ball-to-powder ratio by weight was 10. To minimize oxygen contamination during the milling process, the jars were sealed under a nitrogen atmosphere in a glove box. The milled powder mixtures were granulated through metallic sieves with a #150 mesh size, and then the mixtures were loaded into a graphite mold (inner diameter: 20 mm) in a glove box. The sintering temperature was measured using an optical pyrometer focused on a halfthrough hole of the graphite mold. The inner diameter of the graphite mold was 23 mm and the weight of the powder mixture for making the nozzle neck was 130 g. The powder mixture was densified at 2273 K (heating rate: 100 K/min) by SPS (Dr. Sinter 2020, Sumitomo Coal Mining Co., Tokyo, Japan) for 1 h under a uniaxial pressure of 40 MPa in vacuum (~20 Pa). The bulk density of the sintered specimen was measured using Archimedes' method. The theoretical densities of HfC and SiC are 12.2 g/cm3 and 3.21 g/cm3, respectively, and the values of the composites were calculated according to the rule of mixtures [32]. The ablative graphite was procured from IBIDEN (Ex-60, Japan). The bulk density of the commercial graphite is 1.80 g/cm3. The thermal conductivity and the coefficient of thermal expansion are 110 W/m-K and 5.0 10−6/K, respectively.  2.3. Size limitation  The determination of the insert size needed careful consideration as  Fig. 1. Experimental setup.  33  +=+=HfSi3CHfC2SiCG1673857.456T2×\\x0c\", 'K.-S. Kim, et al.  the HfC showed poor sinterability due to its extremely high melting point of over 3273 K. The sintering and the machining of the HfC-SiC composite into the contoured nozzle shape was challenging due to the super-hard, brittle nature of the ceramic. The HfC-based composite constituted of only a nozzle throat insert, where a severe ablation was expected in a simple tubular shape considering these limitations. The construction of a whole nozzle structure was necessary, including the converging and diverging sections, and all nozzle structures were composed of the ablative graphite, excluding the insert in the case of testing the HfC-SiC ablation. The graphite nozzle was machined precisely in the same dimensions as the HfC-SiC insert and the graphite casing. The HfC-SiC and the graphite insert were embedded within the casing, and the combined nozzles were then mounted onto the rocket system. The nozzle components, including the HfC-SiC composite, are shown in Fig. 2. The characteristic velocity, C*, is not a function of the throat geometry; it exclusively depends on the propellants and the combustion characteristics [29,37]. The tubular throat area, was calculated using the definition of characteristic velocity shown in Eq. (2) and the NASA Chemical Equilibrium with Applications code.  (2)  Here, is the chamber pressure, and is the propellant mass flow rate. The tubular throat diameter and the outer diameter of the nozzle insert was determined as 8.47 mm and 19 mm, respectively.  2.4. Data acquisition & sensor  The pressure of the catalyst bed and chamber were measured with a pressure transducer (PSHD0070BDAJ, Sensys, Korea) with an accuracy of 0.037 % in the gauge pressure. The temperatures of the catalyst bed were obtained using a K-type thermocouple. The authors mounted a K-type thermocouple at the end of the catalyst bed for measuring the temperature of the decomposed HTP gas, as the K-type thermocouple was appropriate for use in an oxidizing environment at temperatures up to 1533 K. A data acquisition system (DAQ) logged the data relating to pressure and temperature. The PXle-4303 and PXle-4353 modules (National Instruments, USA) were connected as a signal conditioning device with a sampling rate of 1 kHz. A dynamic signal amplifier (DN Acta Astronautica 173 (2020) 31-44  AM100, DACELL, Korea) and 100 kgf scale loadcell (CM-100, DACELL, Korea) were used to measure the thrust of the hybrid rocket. The load cell was mounted and aligned using a slide thrust measurement rig.  2.5. SEM morphology & sample preparation  The surface morphologies of the graphite and the HfC-SiC inserts were analyzed by scanning electron microscopy (SEM, Magellan 400, FEI company, USA), with backscattered electron (BSE) and energy dispersive spectrometry (EDS) mapping. After the combustion test, the inserts were mounted in resin and polished with a diamond grind for cross-sectional processing. The cross-sectional image was then investigated to distinguish the oxidized layer from the interior material with BSE and EDS mapping.  2.6. 3D Laser Scanning  A quantitative measurement of the ablated throat was necessary in order to analyze the ablation characteristics. In this study, the measurement of ablation was implemented by adopting a non-contact optical 3D LASER scanning system. The 3D LASER scanning system constitutes two optical cameras and a fringe pattern projector. A 3D structural data is acquired through the triangulation and distortion of the fringe pattern at high resolution. The triangulation calculates the spatial coordinates of the measurement points, and distortions in the fringe pattern reveal the surface boundary. The graphite nozzle and HfC-SiC insert were scanned using 3D Laser Scanning (ATOS III Triple Scan, GOM, Germany). The error in the measurement was estimated as 6.25 μm, based on acceptance testing. The ablated radius and rate were then calculated using the following equations:  (3)  (4)  The ablated radius, was defined as the difference in the nozzle radius before and after combustion and the ablated rate, was defined as the ablated radius divided by the burn time, considering the pressure rising time, . The pressure rising time was marked as the time period from the onset of HTP injection to the stage when the  Fig. 2. (a) Nozzle components and (b) drawing of the Gr-test and the Hf-test.  34  At=ACmPttcPcmt±=aftbfr=bprtbtpr\\x0c', 'K.-S. Kim, et al.  chamber pressure reaches 90% of the initial steady-state pressure. The nozzle surface was considered to not erode simultaneously with the fuel ignition. Instead, we assumed that after the rising time, the chamber pressure successfully develops into the design pressure. Then, significant shear stress and heat flux might be exerted on the nozzle causing ablation. Therefore, the time delay was set as the pressure rising time to ensure that no erosion occurs, and it was subtracted from the burning time, as shown in Eq. (4).  3. Results and discussion  In this study, the ceramic composite was tested under the combustion environment of the hybrid thruster. It was analyzed with comprehensive methods such as quantitative measurement, variation in rocket performance, and investigation of micro-morphologies. The existing ablative graphite was also examined in a carefully controlled condition to distinguish the ablation-resistant mechanism of the HfCSiC composite. The tests with the graphite nozzle and HfC-SiC insert were labeled as the Gr-test and Hf-test, respectively. The temperature trends for both the tests are presented in Fig. 3. The decomposition of the HTP was successfully achieved at temperatures of 1046 K and 1026 K for the Gr-test and Hf-test, respectively, as measured in the catalyst bed. These temperatures are comparable to the adiabatic decomposition temperature of 90 wt % HTP, which is 1022 K. This result indicates the successful decomposition of HTP in both the tests. The decomposition temperature of the Gr-test was slightly higher than the ideal adiabatic temperature; this result is due to the effects of radiation from the combustion gases. Fig. 4 shows an image of the combustion test. The oxidizer-to-fuel ratio is the primary factor affecting the rocket performance. Therefore, we conducted six hot-fire tests, including the Gr-test and the Hf-test, for estimating the fuel mass flow rate by varying the range of the oxidizer mass flux, from 140 to 500 (see Fig. 5). The regression rate of fuel, is mainly dependent on an oxidizer mass flux. The regression rate of fuel, and the mass flow rate of fuel, are obtained using an empirical correlations, which are common practice for estimating fuel mass flow rate in the hybrid motor. The equations are defined as follows.  (5)  (6) where, and are the empirical constants, is the density of fuel, is the fuel length, and is the fuel port diameter. As presented in Fig. 5, the discrete points of time-averaged data were then used to determine the empirical constants; the parameters and were calculated to be 0.0348 and 0.51, respectively. Therefore, the O/F ratio in the initial and the termination phases can be calculated from the fuel burning rate estimated by Eq. (6), and the measured mass flow rate of the oxidizer.  3.1. The 3D-Scanned nozzle contour  The nozzles of both tests were scanned using non-contact 3D LASER scanning, before and after the hot-fire test to obtain the three-dimensional structural data. As shown in Fig. 6, the structural data were processed into 2D contours by slicing 1 mm cross-sections of the nozzle along the direction of flow for comparison. Each cross-sectional area of the nozzle was fitted with an equivalent circle of the same area. The processed 2D nozzle contour and its magnification in the vicinity of the nozzle throat are presented in Fig. 7. The entrance of the converging section and end exit of the diverging section of the nozzles were nearly intact compared to the nominal contours in the Gr-test and Hf-test. Severe ablation was only produced near the nozzle throat. A noticeable erosion on the interfaces (leading and rear boundary) and within the insert was observed in the Hf-test. However, the ablation within the insert decreases while moving  Acta Astronautica 173 (2020) 31-44  downstream. The Gr-test showed smoothly ablated contours along with significant erosion. The minimum radius within the insert for the Grtest and the Hf-test was calculated to be 5.08 mm and 4.62 mm, respectively. The initial throat region, from leading boundary to the rear boundary, has the same smallest cross-sectional area. However, a question may arise regarding the throat location after the irregular deformation of the throat shape. There are two possible ways to determine the throat radius within the insert, which are the minimum insert radius or the averaged radius, as the nozzle contours of both tests are not tubular anymore due to the ablation. The radius of the newly formed throat should be defined. The authors identified the minimum insert area as the determinant parameter for the chamber pressure development. This hypothesis can be supported by the fact that compressible flow is only choked at the throat, as it will consistently accelerate up to the location with the smallest cross-sectional area (i.e., throat). Moreover, the determinant insert radius can be discerned by comparing each reversely-calculated-pressure, , from the minimum and averaged insert radius with the measured pressure, . The can be reversely estimated from the parameters such as the propellant flow rate, the theoretical characteristic velocity, and the throat area using Eq. (2). The values for each test are listed in Table 1. As the realtime throat area cannot be obtained, the characteristic velocity efficiencies, were used as the constant value for the initial stage. The chamber performance parameters ( , and C*) are obtained from Table 2. The specific conditions of calculation and discussion are provided in the next section. The and refer to the reversely-calculated-pressure using the minimum radius and the averaged radius, respectively. The values at the terminal stages of both the tests were close to the measured chamber pressure, whereas the values deviated considerably. The slight difference between the values of and can be attributed to the variation in the value. Moreover, the at the terminal stage of the Hf-test was found to be lesser than that of the Grtest. The calculation of is clearly inconsistent as the chamber pressure during the Gr-test dropped more notably than during the Hftest (see Fig. 13). It is revealed that a chamber pressure developed in order to adjust the newly formed throat, which has the smallest crosssectional area. The minimum insert radius shows adequate physical implications than the averaged radius, although the nozzle contour might irregularly deform due to the ablation. Therefore, the authors defined the newly formed throat as a location that showed the smallest cross-sectional area. This interpretation, again, is in sync with the conventional definition of the throat, which has the minimum area. The  Fig. 3. Temperature trends of the Gr-test and Hf-test.  35  Goxmkg/2rmf=raGoxn=maLmD(4)fnfoxnpn112anfLDpanPrevPcPrevPrevcmtPrevmin,Prevavg,Prevmin,Prevavg,PcPrevmin,cPrevavg,Prevavg,\\x0c', 'K.-S. Kim, et al.  Acta Astronautica 173 (2020) 31-44  Fig. 4. The hot-fire test of the hybrid thruster.  87.51% and 19.19% in the Gr-test and Hf-test, respectively. The HfCSiC clearly showed enhanced ablation resistance for the throat area. A proper interpretation is required to analyze the heat transfer and gas dynamic parameters that are responsible for the ablation. Therefore, a numerical simulation for two-dimensional axisymmetric nozzle flow is carried out using the Ansys fluent 16.1 solver. The compressible flow is based on the shear stress transport (SST) model. The gas mixture composition and the inlet boundary condition (i.e., gas temperature and pressure) are imposed by using the NASA CEA code calculation. The outlet condition is set as the standard atmosphere at the sea level. The wall interface between the combustion gas and solid is specified as a coupled boundary. The thermal and frictional conditions on the wall would be consistently changed as the ablation deforms the throat shape; a numerical investigation of the initial thermal and frictional loads along the tubular nozzle surface might provide complementary information regarding ablation. Fig. 8 and Fig. 9 indicate the relationship of the heat flux and the wall shear stress with the ablation rate for both the tests. The conventional contoured nozzle reveals smooth streams with the isentropic flow; however, flow detachment developed in this case due to the abrupt geometrical change at the leading boundary. The gas separation from the wall induced a significant reduction in the heat flux and the shear stress right behind the leading boundary such that the shear stress dropped approximately to zero. As the operating gas flow repeatedly reattached within the rear tubular throat, the heat flux and the wall shear  Fig. 5. Regression rate of  fuel.  newly formed throat is located near the leading boundary and rear boundary for the Gr-test and the Hf-test respectively as shown in Fig. 7. By comparing each with the initial throat area, the throat is enlarged by  Fig. 6. Processing procedures from the 3D structure data to the 2D nozzle contour: (a) 3d-scanned structure data, (b) structure data with multi-section, and (c) 2D nozzle contour.  36  \\x0c', 'K.-S. Kim, et al.  Acta Astronautica 173 (2020) 31-44  Fig. 7. (a) Nozzle contour of the Gr-test and Hf-test, and (b) the magnified contour along the nozzle throat.  stress rise rapidly and reach a plateau as the flow moves to the rear boundary. This numerical result indicated that a more extreme condition was exerted at the rear part instead of the onset of the insert. The ablation rate of the graphite well resembled the trends of the thermal and frictional loads, with the only exception of the rear boundary. The newly developed throat is located at the position where the stream detached (i.e., minimum thermal and frictional loads), and the throat ablation rate is calculated as 34.09 μm/s. It is found that with an increase in the load exerted on the surface, the graphite erosion increases. However, unlike the graphite, the HfC-SiC composite eroded in a different manner. The most extreme thermal and friction loads are exerted on the HfC-SiC throat, which is located at the rear insert part. It is apparent that the HfC-SiC can withstand ablation significantly even when the throat is exposed to the most severe condition. The throat ablation rate of the HfC-SiC was substantially reduced to 15.81 μm/s, which is 46.5% of the graphite erosion rate. The adoption of the HfCSiC can resist the ablation on the throat under an extreme rocket environment. However, it is also necessary to analyze the occurrence of noticeable ablation at the interfaces and within the HfC-SiC insert. The introduction and the boundaries of the HfC-SiC insert even showed more irregular and severe erosion than the graphite, respectively. The irregular erosion that developed within the front part of the HfC-SiC insert is presumed to be greatly affected by the wall shear stress. The rise and fall of the wall shear stress might affect the erosion of the insert resulting in continuous fluctuations of the ablation rate (see Fig. 9). In addition, peak loads are assumed to be still present in the Hf-test, even after the contour is changed, as the sharp geometrical gradients between the interfaces and the throat paradoxically remain due to the enhanced ablation-resistance of the HfC-SiC. In short, harsher shear stresses are imposed at the interfaces of the Hf-test, leading to severe erosion. On the other hand, the graphite tubular nozzle is considered to be vulnerable to ablation at the interfaces due to the imposition of peak loads. The graphite nozzle deformed its shape to have a smooth curvature at the interfaces and within the insert due to  Table 2 Variation in the chamber performance parameters.  , kg/s  O/F  Theo. Tc, K  Pc, bar  Theo. C*, m/s  Gr-test Initial Stage Terminal Stage Change Hf-test Initial Stage Terminal Stage Change  124.4  146.9  +18.06%  123.5  129.0  9.34  10.24  +0.9  9.32  9.63  2647.77  2566.12  −3.08%  2647.81  2618.71  31.92  25.90  1562.2  1532.4  −18.84%  −1.90%  30.63  27.62  1562.6  1552.1  +4.48%  +0.31  −1.10%  9.82%  0.67%  the ablation, and therefore, a relatively smooth gas flow is developed. Peak loads and fluctuating wall shear stress are no longer exerted on the interfaces and within the insert. Therefore, we presumed that the ablation rate at the boundaries is reduced, and this is the reason for the absence of a peak in the ablation rate at the rear boundary in the Grtest. However, even if harsher conditions are imposed on the HfC-SiC, the reason for the severe erosion of the HfC-SiC at the interfaces and within the insert still remains questionable, if we assume that the HfCSiC composite has superior ablation resistance under a hybrid combustion environment. We reasoned that the fast erosion of the graphite might be the dominant cause of the substantial erosion of the HfC-SiC. The graphite is weak in resisting the thermo-mechanical and chemical erosion effects as compared to the HfC-SiC composite as there is no protection layer generation during the oxidation reaction. Thus, graphite is less ablation-resistant, and it erodes easily. The fast graphite erosion occurred at the interfaces. The HfC-SiC composite, however, is considered to have enhanced ablation resistance due to the presence of refractory oxide; this led to the ablation rate difference between the  Table 1 Reverse calculation of chamber pressure using averaged insert radius and minimum insert radius.  Case  Phase  , g/s  Theo. C*, m/s  Insert radius, mm  Gr-test  Hf-test  Initial stage Terminal stage Initial stage Terminal stage  137.7 161.2 142.4 136.7  1562.2 1532.4 1562.6 1552.1  min.  4.24 5.08 4.24 4.62  5.27  5.10  37  avg.  83.80%  80.98%  , bar  min.  31.92 25.53 30.60 26.67  avg.  23.73  21.88  , bar  31.92 25.90 30.63 27.62  mtcPrevPcmox\\x0c', 'K.-S. Kim, et al.  Fig. 8. Wall heat flux and ablation rate.  Fig. 9. Wall shear stress and ablation rate.  graphite and the HfC-SiC at the interfaces. Therefore, the interior surface of the HfC-SiC is eventually exposed to the combustion gases as the graphite casing that covers the HfC-SiC insert is removed. This interior surface exposure can be a geometrical protrusion producing abrupt flow change along the surface. For this reason, the more the interior surface  Acta Astronautica 173 (2020) 31-44  of the HfC-SiC is exposed to the combustion gases due to the fast graphite erosion, the more ablation would incur on the interfaces. The following visual observations can support the hypothesis that the graphite erosion induces additional erosion on the ceramic. Fig. 10, and Fig. 11 show the nozzle image after the combustion for the Gr-test and the Hf-test, respectively. In the Gr-test, a smooth transition from the convergent section to the throat occurred even while substantial erosion is observed visually on the graphite throat. This geometrical change from tubular to semi-contoured shape appears to be in good agreement with the result reported by Shinn [37]. The crater-like dents are also developed, and this demonstrates possible evidence starting that the graphite chunk is removed from the surface due to the thermomechanical erosion (see Fig. 10). Meanwhile, the interior ceramic surface is clearly found in the region where the graphite casing is eliminated as there is the interfacial difference of the ablation rate, as shown in Fig. 11. This exposure of the interior surface creates an abrupt geometrical change, thereby preventing smooth gas flow. Note that the white oxide layer is removed in the regions where the graphite casing was severely eroded (i.e., substantial exposure of interior surface), whereas the white oxide fragments remained in the region where graphite covered the interior surface of the insert. Moreover, the authors selected two sections (L and S sections) in the same plane, which is perpendicular to the nozzle axis, to compare the additional contribution of graphite erosion on the HfC-SiC ablation qualitatively. The L and S sections represent a significant contrast in terms of occurrence of the fast graphite erosion. Procedures and nozzle contours of the two sections are provided in Fig. 12. The nozzle radius of each section is calculated from the center axis for every 1 mm distance. In the L section, the ablation on the convergent graphite casing occurred at a higher rate than in the HfC-SiC, resulting in the exposure of a large portion of the interior ceramic surface to the combustion gases. Meanwhile, in the S section, it seemed to be that the convergentside graphite eroded at a similar rate with the HfC-SiC; therefore, the interior surface is not exposed. However, this result does not imply that the HfC-SiC has similar ablation resistance as the graphite. It can be explained through the fact that the graphite eroded first, causing exposure of the interior ceramic surface, and then, the ablation rate of the HfC-SiC increased at the leading boundary. When comparing nozzle contours of the S and L sections, the effect of the fast graphite erosion on the HfC-SiC ablation can be clearly found. Significant ablation occurred at the leading/rear boundaries in the L section, whereas alleviated erosion developed in the S section. The nozzle radius increased by 63.44% and 17.92% at the leading boundary for the L and S sections, respectively. For the L section, the interfaces experienced severe ablation, and undesirable irregular contours  Fig. 10. Gr-test nozzle image after combustion.  38  \\x0c', 'K.-S. Kim, et al.  Acta Astronautica 173 (2020) 31-44  Fig. 11. Hf-test nozzle image after combustion.  Fig. 12. Processing of sectional data and HfC-SiC nozzle contours of the selected sections: (a) 3D structure, (b) Slicing the sections, and (c) Contour of the sections.  the ceramic com role in the excess ablation of  have played a major posite. Interestingly, the ablation extent on the L section gradually reduced to the level of the S section while moving back towards the rear boundary. The nozzle radius of both the sections comparably converged into each other, despite the fact that severe interfacial erosion occurred in the L section. The authors postulate that the additional ablation effect due to the fast graphite erosion gradually decreased while moving downstream. Again, the heat flux and the wall shear rapidly increased at the rear part, and the fast graphite erosion induced excessive ablation on the HfC-SiC interfaces. By considering these interpretations, the ablation on the HfC-SiC insert is gradually inhibited while moving downstream due to its enhanced ablation resistance rather than the fact that the insert eroded itself at the front part (see Fig. 7). Our findings conclusively reveal that the adoption of the HfC-SiC into the rocket nozzle insert can significantly alleviate ablation on the throat, even with the highest heat flux and the wall shear stress were exerted. The enhanced ablation resistance of the HfC-SiC insert showed an ablation rate of 46.5% compared to the existing graphite nozzle. Nonetheless, it should be noted that considerable erosion occurred at the interfaces of the graphite casing and the HfC-SiC insert due to the fast graphite erosion. Therefore, it is highly recommended to avoid placing interfaces at the location where substantial ablation would be expected (i.e., vicinity of the throat) for preventing excessive interfacial erosion due to the different ablation rates of the materials.  3.2. Variations in the rocket performance with nozzle ablation  The ablation on the nozzle throat had a significant influence on the rocket performance parameters, including chamber pressure, mass flow rate, and thrust. Each parameter was investigated with a comparison between the Gr-test and the Hf-test.  39  Fig. 13. Chamber pressure and mass flow rate of Gr-test and Hf-test.  developed within the insert. The ablation extent on the entire S section is alleviated considerably than on the L section, even more than the Grtest except for the onset of the divergent section. As discussed previously, the significant wall shear at the rear boundary is present as the difference in the ablation rates causes a sharp geometrical gradient, resulting in an abrupt flow change. Therefore, considerable ablation can still be found at the onset of the divergent side for both sections. The L section showed more significant ablation at the rear boundary than the S section. It is apparent that the fast graphite erosion would  \\x0c', 'K.-S. Kim, et al.  The oxidizer mass flow rate calculated from the MFM differential pressure, and the chamber pressure are presented in Fig. 13. Delays in the ignition observed in the combustion of the fuel grain were 0.340 s and 0.456 s for the Gr-test and the Hf-test, respectively. A surge in the mass flow rate during the transient state is presumed to be because the pressure build-up in the chamber was delayed in the Hf-test, so the over-feeding of the oxidizer occurred. The initial pressure of the Gr-test was measured as 31.92 bar, and this had decreased by 18.84% at termination, whereas the Hf-test showed a 9.82% decrease in the pressure and an initial pressure of 30.63 bar. The drop in the pressure curves for each test were almost linear. The pressure decay slopes, which were linearly fitted after the steady states were achieved, were estimated as −0.28 bar/s and −0.12 bar/s for the Gr-test and the Hf-test, respectively. After developing a steady-state injection in the oxidizer mass flow rate, similar trends were achieved for up to 8 s. However, the final mass flow rate in the Gr-test increased abruptly by approximately 18.06%, while there was only a 4.48% increase in the oxidizer flow rate in the Hf-test. There are three governing factors that can affect the pressure variations, namely, the O/F ratio change, propellant flow rate, and enlargement of the throat area, as shown in Eq. (2). First of all, we evaluated the effect of O/F ratio variation on the chamber pressure, as the characteristic velocity is the direct function of the mixture ratio. Initial O/F ratios are 9.34 and 9.32 for the Gr-test and the Hf-test, respectively. Since the oxidizer mass flow rate for the early stage was carefully controlled, the initial O/F ratios are nearly the same. The oxidizer mass flow rate in the Gr-test increased, as shown in Fig. 13. Therefore, the O/F ratio rose resulting in a ratio of 10.24 at the terminal stage. Slight increase of the O/F ratio in the Hf-test is found to be 9.63. Table 2 summarizes the changes in characteristic velocity, C* depending on the O/F ratio. The theoretical characteristic velocity of each stage is calculated under each O/F ratio and chamber pressure by using the NASA CEA. However, the small rise of the O/F ratio did not affect the characteristic velocity significantly. The theoretical C* dropped only by 1.90% due to an increase in the O/F ratio from 9.34 to 10.24 in the Gr-test, whereas, C* of the Hf-test degraded by −0.67% due to an increase of O/F ratio from 9.32 to 9.63. This result indicates that the variation of the O/F ratio in both tests induced a minor effect on C*. Therefore, the rise of the O/F ratio was considered to have a limited influence on the pressure drop. Secondly, it is apparent that the chamber pressure is the direct function of the throat area, as shown in the following choked flow Eq. (7).  Acta Astronautica 173 (2020) 31-44  the mass flow rate. As the chamber pressure drops due to ablation, the differential pressure between the feeding system and the chamber increases, and this leads to a rise in the mass flow rate. Because no flow control device, such as a cavitating venturi, was mounted in the feeding system to control the flow rate regardless of any changes in the downstream pressure, the chamber pressure and the mass flow rate were coupled. This coupling can explain the increase of the oxidizer flow rate in both tests. A considerable rise of flow rate occurred in the Gr-test; however, the adoption of the HfC-SiC insert can reduce unpredicted and unwanted consumption of the propellant during combustion. An increase in the oxidizer mass flow rate alleviates the drop in the chamber pressure as the flow rate instantly reacts to the downstream pressure variation. This means that the Gr-test showed a more rapid decrease in the pressure even while the propellant is over-fed, which produced an alleviation in the chamber pressure. The tendencies within the pressure curve would presumably be different if flow control devices had been used to maintain the mass flow rate, because the constant flow rate no longer compensates for the decrease in pressure, hence the chamber pressure for the significantly ablated nozzle might drop more notably. If the flow rate is carefully controlled, the pressure of the Gr-test will decrease sharply, while moderate and stable pressure trends would be expected to be developed in the Hf-test. Even when the tubular throat geometry deformed irregularly within the HfC-SiC, the degradation in the chamber pressure is found to be notably alleviated, as compared to the Gr-test. The normalized thrust variations are given in Fig. 14. Each thrust curve was normalized relative to the initial state. In the Hf-test, there was a slight decrease in the thrust, and a subsequent recovery to the initial thrust level. Throughout the Hf-test, the thrust level at the termination phase was nearly the same as the initial thrust, with only a 1% rise. However, the thrust rapidly increased to 26% in the Gr-test. Some of the increase in the thrust is presumed to have been caused by the deformation in the shape from tubular to contoured nozzle, because nozzle discharge and thrust efficiency are generally known to be higher in the contour-shaped nozzles [37]. The rise of the thrust in the Gr-test was primarily produced by the rapid increase in the mass flow rate. Note that the thrust curve was in line with the oxidizer flow rate tendencies (see Fig. 13). In conclusion, it is expected that significant ablation on the rocket nozzle can cause changes in the rocket performance. As the pressure markedly drops in the chamber, the propellant flow rate increases, causing variations in the thrust. The drop in the chamber pressure  (7)  (8)  Here is the specific heat ratio, and is the universal gas constant. As the variations in the temperature and the specific heat ratio were not significant due to the O/F ratio change, the chamber pressure is inversely proportional to the throat area. Therefore, enlargement of the throat area causes the chamber pressure to drop. Although the throat area cannot be measured in real-time, a drop in the chamber pressure occurred to adjust the throat enlargement at the given combustion characteristics. The Gr-test which resulted in an increase of the throat area by 87.51%, showed a more rapid degradation in pressure than the Hf-test, which is in good agreement with the qualitative measurement of the throats. The absolute extent of the pressure drops in the Gr-test is higher than in the Hf-test; the pressure decay slope of the former is substantial than the latter as well. These results indicate that the substantial erosion in the graphite throat severely affects the chamber pressure compared to the HfC-SiC composite. Moreover, we found that the chamber pressure drop is coupled with  Fig. 14. Normalized thrust variation for ablation tests.  40  =PmTAcchokedct=++R1212(1)R\\x0c', \"K.-S. Kim, et al.  causes non-optimal combustion, leading to poor combustion. The flow rate increases, coupled with a drop in the chamber pressure unless the flow rate is carefully controlled, incurring an undesirable and unpredicted consumption of propellant. Moreover, the undesirable increase in thrust leads to a problem with controlling the total impulses of a rocket. Precise control of the rocket may be unattainable due to the significant ablation. The adoption of refractory ceramics may resist this degradation in performance and produce a more stable operation. This result suggests that the insertion of novel refractory ceramics can be verified as feasible for application in rockets, and can alleviate the substantial degradation in performance.  3.3. Material analysis  The density of the sintered HfC-SiC was 11.94 g/cm3, which was higher than the theoretical value of the HfC-SiC composite (11.71 g/ cm3). The high density was attributed to the contamination of WC (theoretical density: 19.59 g/cm3) during milling process. The WC content in the milled HfC-HfSi2-C mixture was reported to increase from 0.259 to 0.363 wt % as the amount of the sintering additives (HfSi2-C) decreased from 30 to 20 wt % because HfSi2 has a lower hardness of 9.95 GPa than the HfC, which has a hardness of 26.5-32.9 GPa [33]. The hardness of the HfC-SiC nano-composites containing 20, 30 and 40 wt % of HfSi2-C additives were measured to be 20.8, 21.2 and 21.5 GPa, respectively [33]. The decrease of hardness with the increase of the additive content was attributed to the low hardness of HfC (18.3-28.4 GPa) compared with that of SiC (26.5-32.9 GPa) [34]. The hardness of HfC-SiC 5 wt % was estimated to be 20.2 GPa by the extrapolation of the hardness values of the composites with 20-40 wt % additives. The Young's modulus of the HfC-SiC nano-composites containing 20-40 wt % of HfSi2-C additives were measured to be 292 GPa. The Young's modulus did not strongly change with the additive content because the Young's modulus of SiC and HfC were similar (299.6-430 vs 424 GPa) [35]. The coefficient of thermal expansion (CTE) of HfC was measured to be 6.87 × 10−6/K at 298-1973 K, which value decreased to 6.51 and 6.31 × 10−6/K when adding 20 and 30 wt % of HfSi2-C additives because the CTE of SiC (4.7 × 10−6/K) is lower than that of HfC. The CTE of the composites with 20 and 30 wt % additives were calculated to be 6.45 and 6.28 × 10−6/K, respectively, by the rule of mixture because the volume of SiC in the composites was 19.3 and 27 vol %, respectively. The calculated CTE values matched well with those of measured ones. The CTE of HfC-SiC with 5 wt % HfSi2-C was estimated to be 6.75 × 10−6/K when considering 5.38 vol % SiC in the composites. Because the thermal conductivity of HfC and β-SiC were reported to be 22.15 and 125 W/mK, respectively, the thermal conductivity of HfC-SiC with 5 wt % additives was estimated to be 27.3 W/mK by rule of mixture [36]. SEM analysis of the polished surface indicated that the porosity of the sintered specimen was less than 2% as presented in Fig. 15. In addition, the formation of fine and homogeneously distributed secondary SiC phase (dark phase) was identified. The density of the sintered specimen was 98% of the theoretical density, and no secondary phases such as residual HfSi2, carbon or oxide phases were detected by XRD. The microscopic surface morphology of the graphite nozzle after the Gr-test is shown in Fig. 16. It was found that a sheet of thin, stacked layers developed uniformly on the surface of the graphite insert. Shattered fragments from the graphite layer were also observed between the layers. This change in structure can be explained by the fact that the graphite surface is oxidized to become a solid phase graphite oxide by the highly oxidizing gases at elevated temperatures. The oxygen that has been permeated into the graphite crystal repulses each layer, so an exfoliation of sheets occurs, forming a graphite oxide [38,39]. This change to a layered surface can cause severe mechanical instability. The mechanically weakened structure is easily damaged because the high-shear combustion gases continuously provide  Acta Astronautica 173 (2020) 31-44  substantial friction to the surface. The existence of the graphite oxide implies that surface erosion occurs in graphite not only by the chemical oxidation, which generates carbon monoxide or dioxide but also via the formation of graphite oxide that weakens the structural stability by allowing oxygen to permeate the graphite crystal. For highly oxidizing rocket environments, the removal of the weakened structure caused by graphite oxide formation could be promoted as it may considerably contribute to the ablation. Therefore, it is suggested to consider that the graphite is eroded from the surface through the combined effect of thermochemical and thermomechanical ablation, which is provoked by the oxidation. The microstructure of the oxidized HfC-SiC is shown in Fig. 17. The “lumpy” surface developed because of the oxidation. EDS elemental analysis confirmed that the surface was constituted of Hf, Si, and O, implying the generation of an oxidized layer. Upon ablation, the fundamental difference between the graphite and the HfC-SiC insert is the existence of the oxide layer. The melting temperature of HfO2, which was formed by the oxidation of HfC, is 3026 K. Because of the excellent thermal stability of HfO2, the oxidation of HfC may be suppressed once the protective oxide layer forms on the surface during the test. The oxide layer formed on the surface in a uniform manner, without any separation between the interior material and the oxide layer. However, the oxide layer contained micro-cracks and pores. The coefficient of thermal expansions (CTE) of HfO2 and HfC were reported to be 5.85 × 10−6 and 6.25 × 10−6/K at 298-1273 K, respectively [35]. The CTE of HfC-SiC with 5 wt % HfSi2-C was estimated to be 6.75 × 10−6/K. Thermal expansion of the oxide itself during the test and the difference of CTE values between HfO2 and HfC possibly increased the residual stress within the ceramic, leading to the development of fractures in the oxide layer. In addition, phase transformation of HfO2 during cooling also induced residual stress. After the test and subsequent cooling down, thermally stable HfO2 phase which formed on the surface of the nozzle throat insert changed sequentially from cubic to tetragonal and monoclinic phase at 2873 and 1993 K, respectively. The phase transformation from tetragonal to monoclinic phase occurs by martensitic mechanism with more than 2% of volume expansion [40]. The micro-cracks propagated easily along the direction with weakened boundaries, providing an increase in the exposure of the interior material to combustion gases. The oxidation of the hafnium carbide also results in the generation of CO gas. Since the dissolubility of CO into oxide scale or silica glass is close to zero, the CO gas escapes the structure, producing submicron scale pores. Increasing porosity on the oxidized surface provides an additional path for the oxygen diffusion [41]. Moreover, the pores aggravate the structural stability of the  Fig. 15. Polished cross-section of HfC-SiC composite showing the dense microstructure.  41  \\x0c\", 'K.-S. Kim, et al.  Acta Astronautica 173 (2020) 31-44  Fig. 16. Graphite surface morphology after combustion.  Fig. 17. Surface morphologies of oxidized HfC-SiC.  oxidized layers, increasing the possibility of cracking. The XRD of the sample after the test could not be measured because of the complex shape. Instead, the XRD data of HfC-SiC composites after oxyacetylene test using a flat sample indicated that only HfO2 phase was detected at the oxidized surface. SiC or SiO2 was not detected at the oxidized surface because of the decomposition of SiO2 and the formation of silicate glass having amorphous structure [28]. The cross-sectional BSE image and EDS mapping of  the HfC-SiC  insert are given in Fig. 18. The original BSE image can be seen in Fig. 18 (a). The formation of pores and cracks in the oxide layer was clearly identified. The rapid gas evolution during the oxidation of HfC, thermal shock caused by the different CTE between the oxide layer and matrix, the volume change during the phase transformation of HfO2 at high temperature as well as the formation of pores and gas during the active oxidation of SiC may be the possible reasons for the formation of cracks in the oxide layer. The rest of the figure is concerned with the EDS  Fig. 18. SEM BSE image and EDS mapping on the cross-section of HfC-SiC composite after combustion (a) BSE image, EDS elemental mapping, (b) Hafnium, (c) Silicon, (d) Oxygen, (e) Carbon, and (f) combined element mapping. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)  42  \\x0c', 'K.-S. Kim, et al.  mapping analysis corresponding to the original BSE image. The upper green layer in Fig. 18 (e) is a resin layer used for adhesion. The crosssectional surface of the insert showed the apparent oxide formation, represented by the grey layer in the BSE image. The oxide layer was mostly composed of HfO2. Some active oxygen diffused into the interior material; however, most of the base materials remained intact even after 25 s of combustion in spite of the formation of cracks and pores in the oxide layer. The thickness of the oxide layer shown as the greyscale in the BSE image (a) was ranging from approximately 63 to 95 μm. The oxidation of HfC occurs rapidly above 973 K with the formation of porous HfO2 layer [42]. However, oxidation was effectively suppressed during the combustion test because the porous HfO2 layer was densified at high temperature and the dense oxide layer effectively suppressed inner diffusion of oxygen during the test (Fig. 18 (a), (d)). In the oxide layer which was mostly composed of HfO2, and SiC, was also oxidized into SiO2, which was evident from the area with Si in Fig. 18 (c) contained oxygen as presented in Fig. 18 (d). The overlapped layer of Hf, Si, and O in Fig. 18 suggests the formation of liquid silicate glass phase composed of HfO2 and SiO2. Also, the eutectic temperature of SiO2HfO2 (1953 K) is lower than the flame temperature [43]. The liquid phase sealed the cracks and pores in the oxide layer and suppressed the inner oxidation of the HfC-SiC composites.  Acta Astronautica 173 (2020) 31-44  problems with precision control in rocket systems. It is therefore suggested that the use of UHTC inserts may have significant advantages for implementation in rockets as compared to ablative materials, by alleviating the variations in performance. 4. The micro-morphologies of the graphite and HfC-SiC specimen revealed the ablation characteristics of both materials. The graphite suffered from oxidation, leading to layered structures and a weakened surface, which was then severely eroded. An oxide layer was found to have formed on the surface of the HfC-SiC insert. This layer was visually observed based on EDM analysis. Pores developed in the surface, which were created during the oxidation of the hafnium carbide. Microscale local cracks were also found. Therefore, a method to prevent the generation of pores and to increase the resistance to thermal shock is necessary for improving the mechanical stability. Possible approaches could be reinforcement with fibers or the use of different ceramics with novel oxidation mechanisms.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  4. Conclusions  Acknowledgements  In this work, the ablation characteristics of graphite and the HfC-SiC composite ceramic were tested with a 250 N scale hybrid thruster using HTP and HDPE. The most important conclusion made may be that it is feasible to use HfC-based refractory ceramic in rocket nozzles, and that UHTCs have inherent advantages in performance compared to ablative graphite. The adoption of UHTCs into the hybrid rocket could possibly overcome the problem with ablation and the associated limitations in operation time although the size of the HfC-SiC composite nozzle throat may be limited in order to minimize the brittle fracture of the ceramic composites. The UHTCs can be considered as promising alternatives for use as a structural material in rocket nozzles. The specific contents of this study can be summarized as follows.  1. The HfC-SiC composite ceramic was used as the rocket nozzle insert and showed an enhanced ablation rate at the throat to the existing ablative material, graphite. The comprehensive analysis for the ablation characteristics of the HfC-SiC insert was conducted based on the rocket performance, non-contact 3D LASER scanning, and the micro-surface morphologies. 2. The structure of the initial and the eroded nozzle specimens were scanned to acquire 3D-structural data. The 3D-structures were then processed into a 2D nozzle contour for comparison. The graphite nozzle had undergone significant erosion, and the nozzle throat had enlarged by approximately 87% compared to the initial state, whereas only 19% of the throat had been enlarged in the HfC-SiC insert. A substantial erosion at the interfaces was found in the Hftest. It was apparent that the fast graphite erosion increases the ablation rate of the HfC-SiC by exposing the interior surface to the combustion gases. This additional ablation contributed by the fast graphite erosion decreased while moving downstream towards the rear part of the insert. Further study is necessary to discern HfC-SiC ablation resistance, excluding the additional contribution of the fast graphite erosion. 3. The chamber pressure and normalized thrust were obtained in both tests. Stable pressure and thrust propensity were observed in the test with the HfC-SiC insert, the Hf-test. In the Gr-test, the pressure reduced severely even though the increased propellant flow rate compensated for the pressure decay. The normalized thrust in the Gr-test increased by approximately 26% due to the increase of the flow rate. 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  "PDF": "Ablation properties of ZrC-SiC-HfB2 ceramic with different amount of carbon fiber under an oxyacetylene flame.pdf",
  "Text": "['Materials Research ExpressACCEPTED MANUSCRIPT • OPEN ACCESSAblation properties of ZrC-SiC-HfB2 ceramic with different amount ofcarbon fiber under an oxyacetylene flameTo cite this article before publication: Maryam Shojaie-ahaabad et al 2020 Mater. Res. Express in press https://doi.org/10.1088/2053-1591/ab70dbManuscript version: Accepted ManuscriptAccepted Manuscript is “the version of the article accepted for publication including all changes made as a result of the peer review process,and which may also include the addition to the article by IOP Publishing of a header, an article ID, a cover sheet and/or an ‘AcceptedManuscript’ watermark, but excluding any other editing, typesetting or other changes made by IOP Publishing and/or its licensors”This Accepted Manuscript is © YEAR The Author(s). Published by IOP Publishing Ltd.(cid:160)As the Version of Record of this article is going to be / has been published on a gold open access basis under a CC BY 3.0 licence, this AcceptedManuscript is available for reuse under a CC BY 3.0 licence immediately.Everyone is permitted to use all or part of the original content in this article, provided that they adhere to all the terms of the licencehttps://creativecommons.org/licences/by/3.0Although reasonable endeavours have been taken to obtain all necessary permissions from third parties to include their copyrighted contentwithin this article, their full citation and copyright line may not be present in this Accepted Manuscript version. Before using any content from thisarticle, please refer to the Version of Record on IOPscience once published for full citation and copyright details, as permissions may be required.All third party content is fully copyright protected and is not published on a gold open access basis under a CC BY licence, unless that isspecifically stated in the figure caption in the Version of Record.View the article online for updates and enhancements.This content was downloaded from IP address 91.243.90.233 on 01/02/2020 at 20:27\\x0c', 'Page 1 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  carbon fiber under an oxyacetylene flame   Maryam Shojaie-bahaabad*, Alireza Hasani-arefi   Faculty of Chemical and Materials Engineering, Sharood University of   Technology, Sharood, Iran, PO Box 3619995161   * E-mail address: mshojaieb@shahroodut.ac.ir   Ablation properties of ZrC-SiC-HfB2 ceramic with different amount of   60 Accepted Manuscript  HfB2 composite was produced by pressureless sintering method. Carbon   milled ZrC-SiC-HfB2. The mixed powders were pressed and sintered at   In the present work, a carbon fiber reinforced ZrC-20 vol% SiC-15 vol%   by oxyacetylene torch. The results showed that as the carbon fiber content   specimens were   investigated by Scanning Electron Microscope (SEM)   fiber with various weight percentages (0, 10, 20, 30) was added to the   toughness tests. The ablation resistance of the composites was performed   equipped with EDS   spectroscopy, XRD   analysis   and hardness   and   2200˚C   for 2h. The microstructure and mechanical properties of   the   Phone and Fax: +982332300253   Abstract   1         \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 2 of 44  surface of the composite underwent a sintering process and formed a dense   after cooling, so these phases played a pinning effect and prevented the   carbon content, the defects in the formed oxide layer on the composite   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  was increased, the porosity and toughness were increased as well, while the   hardness and density were decreased. During the ablation test, the outer   crack propagation. The results also showed that linear and mass ablation   rates were increased with increasing the weight percentage of the carbon   fiber. SEM analysis with EDS further revealed that with increasing the   the pinning effect of ZrSiO4 and HfSiO4 phases as well as the phase   layer of Zr-Hf-Si-O on the surface. ZrSiO4 and HfSiO4 were also created   surface were increased due to the evaporation of SiO2 and the decrease in   60 Accepted Manuscript  fibers are oxidized below 500 °C, Whereas in C/SiC or SiC/SiC the SiO2   transformation of the remained ZrO2 and HfO2. As a result, as the carbon   currently used in aerospace applications. In C/C composites the carbon   application; UHTC; ZrC-SiC.    1. Introduction   C/C and C/SiC composites because of good thermo-mechanical properties,   low thermal expansion, low density and good thermal shock resistance   fiber was increased, the ablation resistance of the composite was decreased.   Keywords:   Composite;   Carbides;   Ablation   properties;   Thermal   2     \\x0c', 'Page 3 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  The coating of C/C composites with a layer of UHTC phase [4, 5].    However, all the mentioned approaches introduce only a little amount of   the UHTC particles in the matrix.    The materials used in temperature protection systems (TPS) need to be   UHTCs are carbides, nitrides and borides of group IV and V transition   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  resistance of these composites as follow:   The reinforcing of C/C composites with UHTC particles.    protective layer prevents the oxidation of SiC matrix up to 1650 °C [1-3].   Several approaches have been used to increased ablation and oxidation   growing demand for advanced structural materials with sustainability at   stable at high   temperature oxidation (>2000˚C). So,   there has been a   60 Accepted Manuscript  conductivity (20.5 Vm-1K), high electrical conductivity (78 × 10-6 Ω cm)   metals such as ZrB2, HfB2, ZrC, TaC, HfC, HfN and ZrN [9]. Zirconium   its high melting point (3420˚C), it has been placed in the UHTCs class [10].   ZrC, like other compounds of this class, has interesting properties such as   hyper sonic aerospace vehicles have attracted much attention   [7, 8].    high corrosion resistance  to acidic and alkaline environments, low thermal   high   temperatures   [6]. Recently,   the ultra-high   temperature ceramics   carbide (ZrC) is an extremely hard and refractory ceramic material due to   (UHTCs) used in temperature protection systems and other components of   3     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 4 of 44  due to the presence of metal bands, high hardness (25 GPa) owing to the   cracked layer is formed on the surface [14]. One way to improve the   4   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  ZrC oxidation begins at the low temperature of about 600 ˚C, depending on   the partial pressure of oxygen. A dense   layer of C-ZrO2   is   formed   preventing further carbon oxidation in the ZrC matrix. At temperatures   temperatures   (1900-2500˚C)   [11-13]. The main drawback of all   the   lower density, as compared with other transition metal carbides like WC,   temperatures higher than 3000˚C and suitable dimensional stability at high   transition metal borides and carbides, are their poor oxidation resistance.   above 800 ˚C, when the partial pressure of oxygen is increased, ZrC is   TaC or HfC, mechanical strength   in corrosive environments even at   presence of strong ZrC covalent bonds, great young module (440 GPa),   60 Accepted Manuscript  through the formation of the B2O3 layer filled with ZrO2 or the formation of   rapidly oxidized to monoclinic or tetragonal ZrO2. As a result, a porous and   a SiO2 or ZrSiO4 protective layer [15]. Researches have also shown that by   additives such as ZrB2, SiC or silicids to prevent the diffusion of oxygen   increasing SiC   the   strength and   toughness of ZrC ceramics can be   oxidation and ablation resistance of ZrC is to prepare composites with   modulus [15]. Adding a higher capacity cation is another way to improve   improved; however,   it does not   influence   the hardness and elasticity   the oxidation and ablation resistance of ZrC ceramics. The addition of     \\x0c', '1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  [16]. Some works have been done on UHTCs reinforced with carbon   at 1500 °C [24]. Recent studies on the kinetics of oxidation of a carbon   Page 5 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  higher capacity cations causes the oxygen vacancies in ZrO2 to be filled,   Also, the addition of higher capacity cations increases the viscosity of the   glass phase, thereby raising the glass boiling point [10].   spheres or short fibers [17-21]. There are few works concerning   the   performance of UHTC-rich matrix under oxidizing atmospheres. Sciti and   Another drawbacks of UHTCs are their low fracture toughness and high   density. Simple method is proposed to fabricate UHTC based composites   Zoli have successfully obtained carbon fiber reinforced UHTCs via slurry   with short carbon fibers as both a reinforcing and mass-lightening phase   resulting in the stability of the ZrO2oxide layer in the tetragonal structure.   60 Accepted Manuscript  based on ZrB2 with 0-50 vol% of short carbon fibers at 2500 °C under   fibre reinforced ZrB2 composite doped with 10 vol.% SiC have shown that   oxyacetylene   torch for 60 s [16]. Elsewhere, ZrB2/SiBCN composites   containing carbon fibers coated with BN were prepared using combination   infiltration and hot pressing [22, 23]. Preliminary studies on these materials   have shown their promising mechanical properties and oxidation resistance   the critical   temperature for   these composites   is below 1000 °C [25].   Silvestroni studied ablation properties of functionally graded composite   of sol-gel and spark plasma sintering (SPS) techniques by Miaoa et al.   5     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 6 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  discontinuous SiO2-rich   oxidation treatment [28].     Also.   it has been shown   the oxidation behavior of   the ZrB2-20vol%SiC-2vol%B4C   composite   architecture of   the carbon   fibers on   the oxidation behaviors of   the   Theses researchers compared ablation behavior of ZrB2/SiBCN with and   composites by oxidation tests by an oxyhydrogen burner at 1700°C for 10   the constituent materials [27]. The effect of short carbon fiber addition on   studied by Das et al. Oxidation study was carried out at 1600oC for 2 h.   After the oxidation test, the carbon fiber reinforced composite showed   min. The experimental results showed that the oxidation of both composites   significantly depended on the microstructural parameters and properties of   without Cfs for 10−30 s under oxyacetylene torch [26]. Inoue et al. studied   formation of SiO2 rich, continuous and protective top layer of about 30 µm.   60 Accepted Manuscript  fiber reinforced ZrC-20 vol% SiC-15 vol% HfB2 composite with different   composites reinforced with carbon fibers. In this study, therefore, a carbon   weight percentages of fiber, from 0%   to 30%, was fabricated by   the   resistance and melting point than ZrB2 and ZrO2.    On   the other hand,   the base composite   showed   the   formation of a   that HfB2 and HfO2 have higher oxidation   top   layer of about 20-30 µm after   the same   the ablation behavior of ZrC-SiC-HfB2    There   is still no   literature on   6     \\x0c', 'Page 7 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  pressureless sintering method (PLS). The main purpose of this paper was,   in fact, to study the ablation resistance of the prepared composite.   2. Material and methods    1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  diameter and the average length of 1mm) were used as the reinforcement.   The SEM images of the raw materials used in this study are shown in   Figure 1. Initially, the raw materials were milled by planetary ball mill   using zirconium balls and ethanol. 1 ml of ethanol was added for every 5 g   composite was prepared by the pressureless sintering method. Accordingly,   ZrC powder with the particle size less than 10 µm was used as the matrix.   SiC (< 4µm), HfB2 (< 5µm), and chapped carbon fibers (with 1.5 µm   In this research, a carbon fiber reinforced ZrC-20 vol% SiC-15 vol% HfB2   60 Accepted Manuscript  specimens were then compressed at a pressure of 200 MPa using CIP to   wt%) were then added to the ZrC-20 vol% SiC-15 vol% HfB2 slurry, and   again ball milled for 3 minutes to obtain a homogeneous mixture. The   increase the green body strength. The pressureless sintering process was   also carried out in a high-temperature graphite furnace at 2200 °C under an   obtained slurry was then dried in an oven. After that, the mixture was   compressed at a pressure of 30 MPa using   the uniaxial press. The   of powders. The ball to powder weight ratio was 10:1 and the speed of ball   milling was set at 200 rpm for 3 h. The carbon fibers (0, 10, 20 and 30   7     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 8 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  ZSH20 and ZSH30.   argon atmosphere for 2 h with heating and cooling rates of 10°C/min. In   the present work, the prepared composites were labeled as ZSH0, ZSH10,   60 Accepted Manuscript  Figure 1. SEM images of (a) ZrC, (b) SiC and (c) HfB2 (d) carbon fibers   resistance of the specimens. The oxyacetylene flame was consisted of a   used in the present research.   The oxyacetylene   ablation   test was used   to determine   the   ablation   2.1.   Ablation test   8       \\x0c', 'Page 9 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  ablation test.     9   The test time was 60 s and the surface temperature of the samples was   measured using an   infrared   thermometer. After   the ablation   test,   the   samples were cooled at room temperature. The ablation efficiency was   mixture of oxygen and acetylene gases with a pressure of 1.6 bar and 2 bar,   ablation test.   were as follows:   measured based on the linear ablation rate and the mass ablation rate. The   mass and thickness of the samples were determined before and after the   exposed to the flame up to 3000°C at a distance of 10 mm of the nozzle tip.   The formulas of the mass ablation rate (Rm) and the linear ablation rate (Rl)   and 2 Nm3/h, respectively. The surface of the specimens was vertically   respectively; also the oxygen and acetylene gas flow rates were 1 Nm3/h   60 Accepted Manuscript  Where m0 and l0 are the mass and thickness of the samples before the   ablation test, respectively, and mt and lt are the mass and thickness of the   samples after   the ablation test, respectively; also, t is the time of   the                                   \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 10 of 44  2.2.   characterization   Archimedes principle was used. The formulas used were as follows:   and microstructure of the samples before and after the ablation test. The   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  Where P is the open porosity of the sample, D is the density of the sample,   results of the three samples were selected as the final result. According to   calculated.   In order   to calculate   the   theoretical density values of   the   In order to obtain the bulk density and the amount of open porosity, the   the law of mixtures, theoretical density values for the composites was   M1 is the weight of the dried sample, M2 is the weight of the immersed   sample in water and M3 is the weight of the wet sample. The average   60 Accepted Manuscript  g/cm3 were used as the density of ZrC, SiC, HfB2 and carbon fiber (Cf),   composites,  the values of 6.73 g/cm3, 3.22 g/cm3, 10.2 g/cm3, and 1.8   according to the ASTM 0327-08 .The standard applied load was 2 kg with   respectively. The crystalline phases were identified by X-ray diffraction   dispersive spectroscopy (EDS) was also used to investigate the morphology   hardness of   the   composites was measured by   the Vickers method,   (XRD). A scanning electron microscope (SEM) equipped with energy   the holding time of 10 s on the polished surface. The fracture toughness   10                                       \\x0c', 'Page 11 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  of 15 s.    3. Results and discussion   (KIC) was measured by the direct crack measurement method (DCM), using   in figure 2, the particle size of the powder decreased after ball milling for   minutes, the fiber length was decreased by 200 μm, but the fibers were not   powders used for composite preparation after ball milling for 3h. As shown   Figure 2 shows the scanning electron microscope image of the mixed   3h. Also, after the addition of chapped carbon fibers and ball milling for 3   the Vickers indentor and the applied load was 30 kg with the holding time   60 Accepted Manuscript  uniformly dispersing short fibers in ceramic slurry, although the mixing   powders. The absence of the clumps fibers suggested that the short fibers   time was very short (only 3 minutes). Short mixing time greatly reduced   were uniformly dispersed   in   the matric slurry after ball milling. This   indicated   that   the mixing method used   in   this study was effective   in   dispersed   in   the powders without any clumps   in   the resulting mixed   randomly and uniformly   the damage to the fibers.   damaged.    It was   found   that   the chapped   fibers were   11     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 12 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  milling.   12   60 Accepted Manuscript  Figure 2. SEM images of (a) ZSH0, (b) ZSH10, (c) ZSH20, (d) ZSH30, (e)   and (f) high magnification of (c) and (d) of mixed powders after ball       \\x0c', 'Page 13 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Figure 3 shows   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  the XRD patterns of   the prepared composites. In all   additional phases were formed during the sintering process.    samples there were peaks associated with ZrC, SiC and HfB2 phases. No   60 Accepted Manuscript  phase was ZrC, and the light gray phase and the dark gray phase were HfB2   composites are shown. According to XRD and EDS, it shown that the light   Figure 3. X-ray diffraction of (a) ZSH0, (b) ZSH10, (c) ZSH20 and (d)   Figure 4 shows   the SEM   images of   the polished cross-section of   the   and SiC, respectively. The   isolated pores were formed and uniformly   ZSH30 composites.   13       \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 14 of 44  fibers in the ZrC matrix.   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  distributed in the ZrC matrix. EDS results revealed the presence of chapped   60 Accepted Manuscript  Figure 4. SEM images of (a) ZSH0, (b) ZSH10, (c) ZSH20 and (d) ZSH30   composites after sintering.   14       \\x0c', 'Page 15 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  [29].    15   during the pressing process [21].   In Table 1, the measured densities and mechanical properties of the dense   composites made by the pressureless sintering method are presented. As   The reduction in the fiber length can be explained as follows:    less than that of the primary fibers. In other words, it indicated that the   length of the fibers was changed after the ball milling and pressing process.   First, the fibers were broken during the ball milling process. Second, some   fibers were broken due to the volumetric shrinkage of the green body   On the other hand, with the increase of carbon fiber in each composition,   agglomeration of fiber increases (Figure 5) and this prevents complete   the length of the fibers on the polished cross-section of the composite was   sinterability of composites. According to Figure 5c, it could be seen that   60 Accepted Manuscript  coefficient (CTE) between the carbon fiber, SiC and HfB2 and ZrC matrix   The mismatch between the matrice and the reinforcement phase caused   stresses caused the separation of the matrice/reinforcement interface, as a   result, porosity was created. On the other hand, as the carbon fiber was   tensile and compressive stresses in the matrice and the second phase. These   can be seen, by increasing the carbon fiber, the porosity was increased. The   porosities were   derived   from   the   difference   of   thermal   expansion     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 16 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  increased in the composite, agglomeration was also increased (Figure 5); as   a result, complete sintering of the composites was prevented.   60 Accepted Manuscript  Density was also decreased with the increase in the carbon fiber content.   Figure 5. Agglomeration of carbon fibers in (a) ZSH10, (b) ZSH20 and (c)   Two main reasons for the decrease in relative density are as follows:   ZSH30 composites.   16         \\x0c', 'Page 17 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  observed in other research [29].   1) Agglomeration of primary powder particles during ball milling   accelerated coarse aggregation before densification was complete.   Table 1. The compositions, densities and mechanical properties of the   amounts of the carbon fiber was used, the grain growth was increased.   2)  In the presence of oxide impurities, densification was prevented.   When the carbon fibers were agglomerated and separated at the grain   As can be seen, slight changes in the hardness were seen as a function of   the amount of   the short carbon fiber. A similar result has also been   words, the diffusion at the ZrC grains was increased. Therefore, when large   boundaries, the diffusion at the grain boundaries was decreased; in other   60 Accepted Manuscript  Sample   Theoretical   Relative   Bulk   density   density   density   porosity   (GPa)   Open   Hardness   Toughness   (g/cm3)   (g/cm3)   4.2   4.6   4.7   4.9   91.4   4.91   89.5   4.11   87.2   3.59   SZH20   4.59   SZH10   5.37   SZH30   4.12   composites.   MPa m1/2   3.8   5   7.6   8   (g/cm3)   12.3   11.2   10.4   95.6   6.4   SZH0   6.69   (%)   8.6   17       \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 18 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  (˚C)   18   improvement was related to the deflection or bridging of the crack.    Furthermore, sintered composites showed   that fracture   toughness was   through the surface [30]. The linear and mass ablation rates of the samples   ablation. The rate of ablation could be controlled by the oxidation reactions   increased, the composites showed higher linear and mass dissipation rates   showed possess similar variations [39]. As the carbon fiber content was   when the amount of the fiber was varied from 0 to 30%. Obviously, this   Table 2 presents the linear and mass ablation rates of the samples af ter   improved due to the use of the carbon fibers in the composite matrix (Table   of C [31-32], ZrC [32], SiC [33-37], HfB2 [38] and the diffusion of oxygen   1). Composite fracture toughness changed in the range of 4.2-5.4 MPa m1/2,   60 Accepted Manuscript  Table 2. Linear and mass ablation rates of the composites after ablation.   due to the higher percentage of carbon fiber and porosity [40-45].    Sample   Ablation Time   Ablation   1.64 × 10-4   1.73 × 10-4   2.49 × 10-3   1.62 × 10-3   3.51 × 10-3   2.46 × 10-3   2873   2826   2822   2810   1.05 × 10-3   1.22 × 10-3   Liner   Mass Ablation   (s)   surface   Ablation   (g/s)   temperature   (mm/s)   60   60   60   60   SZH0   SZH10   SZH20   SZH30     \\x0c', 'Page 19 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  and burning carbon fiber played an   important role   in determining   the    weight loss.   Figure 6 shows the XRD patterns of the composite after ablation. It can be   HfO2. The melting point of ZrO2 and HfO2   is 2810 and 2800   ˚C,   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  As the amount of the carbon fiber was increased, the porosity in   the   composite was increased, so the channels of oxygen penetration into the   matrix were increased. As a result, the endothermic oxidation reactions   were accelerated and the surface temperature was decreased after ablation.   composite weight was expected. However, the severe evaporation of SiO2   seen that all patterns were similar and the major oxides were ZrO2 and   In the first stage, ZrC and HfB2 were oxidized; therefore, an increase  in the    60 Accepted Manuscript  respectively [18]. It indicates that the dissipation of ZrO2 and HfO2 was   The melting point of SiO2 is about 1673˚C, showing that SiO2 was melted   At the same time, in the XRD pattern, some diffraction peaks of HfSiO4   room temperature after ablation and prevented the crystallization of SiO2.   and ZrSiO4 could also be observed [46]. In the XRD patterns of ZSH0, no   and evaporated during ablation. Also, scouring flame was high which could   easily spread SiO2 around. In addition, the samples were cooled rapidly at   In the XRD patterns, no diffraction peaks related to SiO2 were observed.   low during ablation.    19     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 20 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  carbon peak was observed, which means that the surface of the sample was   well covered by ZrO2 coating and no carbon material was burnt [39].   60 Accepted Manuscript  sintering process, a dense layer was formed on the surface which could   The SEM images of the composites surface after ablation are shown in   Figure 7-10. It can be seen that the grains were interconnected. The stacked   oxide particles began to sinter at the ablation temperature [47]. During the   Figure 6. XRD patterns of (a) ZSH0, (b) ZSH10, (c) ZSH20 and (d)   ZSH30 composites after ablation.   Three regions could be found on the ablated surface of the composites:   center region, transition region and brim region.    20       \\x0c', 'Page 21 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  prevent oxygen penetration. In addition, this structure was derived from the   high melting points of HfO2 and ZrO2 and   their high viscosity and   these prevented to help the surface from burning and oxidizing [48].   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  resistivity to the spread particles at the high flux of the oxyacetylene flame;   60 Accepted Manuscript  Figure 7. Surface morphologies and EDS analysis of ZSH0 composite in   the ablation (a) center, (b) transition and (c) brim region after ablation.   21       \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 22 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 8. Surface morphologies and EDS analysis of ZSH10 composite in   the ablation (a) center, (b) transition and (c) brim region after ablation   22       \\x0c', 'Page 23 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 9. Surface morphologies and EDS analysis of ZSH20 composite in   the ablation (a) center, (b) transition and (c) brim region after ablation.   23         \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 24 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 10. Surface morphologies and EDS analysis of ZSH30 composite in   the ablation (a) center, (b) transition and (c) brim region after ablation.   24       \\x0c', 'Page 25 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  formed during the ablation process [49] (Figure 7). EDS analysis (Figure 7 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  phases; so, these could play the pinning effect to prevent cracking [50].   The formation of the porous oxide layer was related to the generation of   However, there were still some pinholes distributed on the surface. These   pinholes were generated by the evaporation of gases generated by oxygen   diffusion channels. In the sample ZSH0, no micro-cracks were observed on   10) also showed that the amount of Si was lower than the initial amount in   the precursor powder, indicating that some SiO2 formed during the ablation   could be reacted with ZrO2 and HfO2, generating HfSiO4 and ZrSiO4 stable   the surface; this could be related to the new HfSiO4 and ZrSiO4 phases   process had been evaporated. In addition, some of the remaining SiO2   60 Accepted Manuscript  in the generated HfSiO4 and ZrSiO4, resulting in the weakening of the   SiO, CO2, CO and B2O3 gases, as well as the active oxidation of SiC and   porosity and consequently, oxygen penetration into the matrix, leading to   the production of more gaseous products. According to the EDS analysis of   Increasing   the amount of carbon fiber   in   the composite   increased   the   the ablated carbon fiber in the composite, this may lead to further reduction   the carbon fiber.   pinning effect.   25     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 26 of 44  In   addition,   the   remaining ZrO2   and HfO2   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  can   undergo   a   phase   porosities formed in the oxide layer [6].   decreased. This could encourage the active oxidation of SiC and the coarser   In general,   these defects became   transition channels for   the oxidizing   transformation during   the cooling process   leading   to   the formation of   appear due to the CTE mismatch of the ablation layer and the composite   matrix [53]. These resulted in increasing the porosities and cracks in the   With increasing the porosity in the composite and enhancing oxidation   reactions, the partial pressure of oxygen in the ablation environment was   cracks due to the volumetric effect [51-53]. Furthermore, cracks could   oxide layer formed on the composite surface during the ablation process.   60 Accepted Manuscript  temperature of HfO2 and ZrO2, the SiO2 product in HfO2 and ZrO2 could   cause the lower melting point of HfO2 and ZrO2. It resulted in a continuous   field consisting of oxide (HfO2 and ZrO2) and carbide (ZrC) compounds in   percentages of the carbon fiber in the three ablation zones are shown in   SEM images of the ablated cross-section of the composite with different   Although   the ablation surface   temperature   is   lower   than   the melting   which HfO2 and ZrO2 grains were scattered (Figure 6).   species thereby increasing the composite ablation.   Figure 11-14.   26     \\x0c', 'Page 27 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 11. Cross-section micrographs of ZSH0 composite in the ablation   (a) center, (b) transition and (c) brim region after ablation.   27                   \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 28 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 12. Cross-section micrographs of ZSH10 composite in the ablation   (a) center, (b) transition and (c) brim region after ablation.   28                       \\x0c', 'Page 29 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 13. Cross-section micrographs of ZSH20 composite in the ablation   (a) center, (b) transition and (c) brim region after ablation.   29                     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 30 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  60 Accepted Manuscript  Figure 14. Cross-section micrographs of ZSH30 composite in the ablation   (a) center, (b) transition and (c) brim region after ablation.   30                               \\x0c', 'Page 31 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  formed after the ablation.   31   was sintered and some porosity and micro-cracks were observed.   the pinning effect of HfSiO4 and ZrSiO4 phases.   The dense structure  also revealed that  the oxide layer containing oxide   and carbide particles effectively acted  as a layer against heat and prevented    further oxidation of the matrix.   SEM images of the ablation cross-section of the sample ZSH20 (Figure 13)   showed  that the oxide grains were stacked together with clear boundaries   SEM images of the ablation cross-section of the sample ZSH10 (Figure 12)   showed that due to the relatively high surface temperature, the oxide layer   These defects were caused by the evaporation of SiO2 and the decrease in   60 Accepted Manuscript  represented a coral-like oxide structure consisting of ZrO2, HfO2 and ZrC   evaporation of SiO2 created channels for the penetration of oxygen into the   matrix, which could not be beneficial to prevent the passage of oxygen.   and the structure of the stacked grains was loose. The cracks and pinholes   in the oxide layer were generated by the gaseous products, and the severe   According to Figure 14, the oxidation layer was loose and brittle with a   number of pores and spalls. Thus, one would expect a lack of proper   The cross-section of the sample ZSH30 oxide layer, as shown in Figure 14,     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 32 of 44                                                              (8)                                                  32   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  three processes:   3Evaporation of SiO2   occur:   adhesion between   the oxide   layer and   the composite, so   it could not   1Oxidation of matrix, ZrC, SiC, HfB2 and carbon fiber   2Oxygen penetration through oxide species   effectively prevent base oxidation.   3.1.   Ablation Mechanism   The ablative behavior of the composite can be described by the following   Depending on the XRD patterns, the following chemical reactions might   60 Accepted Manuscript                                                      (2)                                                   (3)                                                                      (4)                                                                     (5)                                                     (1)                                                                          (6)                                                                        (7)     \\x0c', 'Page 33 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1                        (s)                                      (9)   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  Zr-Si-Hf-O layer on the surface. A massive amount of gaseous products   Reactions (1-10) can occur in all regions of the ablation surface, and the   reactions (11 and 12) may occur in high temperature regions, especially in   the surface ablation layer [54].                                                                                 (12)                                                                (10)                                                                            (11)   When ablation begun, the carbides, HfB2 and C were oxidized, forming a   60 Accepted Manuscript  products that were rapidly evaporated to CO (g) and CO2 (g) (reactions 6   was  developed by carbon fiber, ZrC, HfB2 and SiC oxidation reactions   (reactions 1-5). Also, SiO2 (l) and B2O3 (l) were transformed to gaseous   higher evaporation of the  gaseous products, resulting in the increased   composite porosity, more oxygen penetration  to the base;  then, with the   Release gaseous products at all stages of ablation represent an endothermic   process absorbing heat. This could result in lower surface temperatures   and 7). In addition, the reaction (10) also occurred at the high temperature   ablation.   [55]. The higher mass and linear ablation rate of the composite caused the    33     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 34 of 44  increase in carbon fiber, oxidation reactions are accelerated [55]. Due to the   products, a porous oxide species was formed on the surface.   accelerate the sintering process [55,56].   increase of the  carbon fiber in the composite, oxide species were formed   oxidation of carbon fibers and   the matrix and evaporation of gaseous   process, which could prevent the propagation of cracks. However, with the   with further and larger pinholes and a large number of micro-cracks on the   continuous solid solution with ZrO2 in the heating process, so it could   In addition, SiO2 could react with HfO2and ZrO2 (reactions 8 and 9),   forming  the  new stable phases HfSiO4 and ZrSiO4 during the cooling   According to the HfO2-ZrO2 binary phase diagram, HfO2 could form a   60 Accepted Manuscript  ZrC-SiC-HfB2 composites with different weight percentages of carbon   surface due to the  phase transition of HfO2 and ZrO2 and further reduction   amount of carbon   in   the composite was   increased,   the porosity and   oxyacetylene torch for 60s. The rates of linear and mass ablation were   toughness were improved, while hardness and density were declined. The   ablation   characteristic   of   the   composites was   investigated with   an   of SiO2.   4. Conclusions    fiber were prepared by free pressureless sintering at 2200˚C for 2h. As the   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  34     \\x0c', 'Page 35 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  raised with   increasing   the carbon content. Structural evaluation of   the   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  composite were increased and the ablation resistance of the composite was   decreased with increasing the carbon fiber content.   element was decreased. Thus, by reducing   the amount of ZrSiO4 and   content, the defects were increased on the oxide layer surface and the Si   layer was   increased. As a result, oxygen penetration channels   to   the   formed oxide layer also showed that a dense layer was formed on the   HfSiO4 phases, due to the locking effect and the phase transformations of   phases of ZrSiO4 and HfSiO4 were formed. With increasing the carbon   the remained ZrO2 and HfO2, the crack formation in the Zr-Hf-Si-O oxide   composite surface. Due to the reaction of SiO2 with ZrO2 and HfO2, stable   60 Accepted Manuscript  fluxes of C/C composites modified by ZrB2-ZrC and ZrB2-ZrC-SiC   [2] L. Liu, H. Li, W. Feng, X. Shi, K. Li, L. Guo, Ablation in different heat   oxyacetylene torch testing and characterisation, J. Eur. Ceram. Soc. 33   [1] A. Paul, S. Venugopal, J.G.P. Binner, B. Vaidhyanathan, A.C.J.   Heaton,   P.M. Brown, UHTC-carbon   fibre   composites:   preparation,   References   (2013) 423-432.   particles, Corros. Sci. 74 (2013) 159-167.   35     \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 36 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  https://doi.org/10.1016/j.ceramint.2018.02.093.   Soc. 32 (2012) 947-954.   36   91 (2008) 1495-1502.   Self-healing ZrB2-SiO2   flame, Int. J.   Antimicrob. Agents. 6 (2009) 145-150.   oxidation   resistance   coating   for SiC   coated   carbon/carbon composites, Corros. Sci. 110 (2016) 265-272.   [3] E.L. Corral, R.E. Loehman, Ultra-high-temperature ceramic coatings   [4] O. Haibo, L. Cuiyan, H. Jianfeng, C. Liyun, F. Jie, L. Jing, X. Zhanwei,   [5] H. Li, L. Zhang, L. Cheng, Y. Wang, Ablation resistance of different   for oxidation protection of carbon-carbon composites, J. Am. Ceram. Soc.   [6] X. Chen, Q. Feng, H. Zhou, Ablation behavior of three-dimensional Cf   coating structures for C/ZrB2 -SiC composites under oxyacetylene torch   60 Accepted Manuscript  [7] T. Liu, Y. Niu, C. Li, J. Zhao, Effect of MoSi2 addition on ablation   [8] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Oxidation of ZrC-30 vol%   SiC composite in air from low to ultrahigh Temperature, J. Euro. Ceram.   reactive melt infiltration, https://doi.org/10.1016/j.corsci.2018.02.011.   /SiC-ZrC-ZrB2 composites prepared by a   behavior   of   ZrC   coating   fabricated   by   vacuum   plasma   spray,   joint process of sol-gel and     \\x0c', 'Page 37 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  10.1016/j.ijrmhm.2017.03.019.   37   ceramics route, Ceram. Int. 40 (2014) 5025-5031.   temperature oxidation, Mater. Chem. Phys. 143 (2013) 407-415.   [9] D. Pizona, L. Charpentierb, R. Lucasa, Oxidation behavior of spark   resistance of C/C-ZrC-SiC Composites, Mater. Desi. 129 (2017) 15-25.   [12] B. Núñez-Gonzáleza, A. Ortiza, F. Guiberteaua, Spark-plasma [11] Y. Jia, X. Yao, J. Sun, H. Li, Effect of ZrC particle size on the ablation   with   tantalum silicide: Microstructure, mechanical properties and high   [10] L. Silvestroni, D. Sciti, M. Balat-Pichelin, Zirconium carbide doped   plasma sintered ZrC-SiC composites obtained from the polymer-derived   sintering kinetics of ZrC-SiC powder mixtures subjected to high-energy   60 Accepted Manuscript  [13] Z. Nasiri, M. Mashhadi, Microstructure and mechanical behavior of   [14] D. Yung, B. Maaten, M. Antonov, I. Hussainova. Oxidation of spark   plasma   sintered   ZrC-Mo   and   ZrC-TiC   Composites,   https://doi:   co-ball-milling, Ceram. Int. 39 (2013) 9691-9697.   ternary   phase   ZrB2-SiC-AlN   nanocomposite,   https://doi:10.1016/j.ijrmhm.2018.09.009.       \\x0c', 'AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  Page 38 of 44  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  1450°C, J. Adv. Ceram. 6 (2017) 110-119.   Ceram. Int. 39 (2013) 1567-1574.   38   [16] L. Silvestroni, C. Melandri, V. Venkatachalamb, J. Binner, Merging   183 (2019) 1078-1080.   (2013) 413-421.   [15] G. Antou, M. Ohin, R. Lucas, Thermomechanical properties of a spark   plasma sintered ZrC-SiC composite obtained by a precursor derived   ceramic route, Mater. Sci. Eng. A. 643 (2015) 1-11.   [18] W. Hong, K. Gui, P. Hu, X. Zhang, S. Dong, Preparation and   [17] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, The effect of a graphite   addition on oxidation of ZrB2-SiC in air at 1500°C, J. Eur. Ceram. Soc. 33   characterization of high-performance ZrB2-SiC-Cf composites sintered at   toughness and oxidation resistance in a light ZrB2 composite, Mate. Des.   60 Accepted Manuscript   [19] C. Piriou, O. Rapaud, Sintering and oxidation behavior of HfB2-SiC   of short pitch-based carbon fiber-reinforced HfB2-SiC matrix composites,   composites   from 0   to 30 vol.% SiC between 1450   and 1800 K,   https://doi.org/10.1016/j.ceramint.2018.10.075.   [20] S. Guoan, K. Naitoa, Y. Kagawa, Mechanical and physical behavior s     \\x0c', 'Page 39 of 44  AUTHOR SUBMITTED MANUSCRIPT MRX2-103688.R1  http://dx.doi.org/10.1016/j.matdes.2018.02.061.   (2017) 129-138.   behavior, J. Euro. Ceram. Soc. xxx (xxxx) xxx-xxx   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59  [21] F. Yang, X. Zhang, J. Han, Processing and mechanical properties of   Des. 29  (2008) 1817-1820.   213.   and   regression   tree   analysis,   Mater.   Des.   from oxidation in novel UHTCMC materials, Mater. Des. 113 (2017) 207-  [23] D. Sciti, A. Natali Murri, V. Medri, L. Zoli, Continuous C fibre   [24] A. Vinci, L. Zoli, D. Sciti, C. Melandri, S. Guicciardi, Understanding   the mechanical properties of novel {UHTCMCs} through random forest   short carbon fibers toughened zirconium diboride-based ceramics, Mater.   composites with a porous ZrB2 Matrix, Mater. Des. 85 (2015) 127-134.   [22] L. Zoli, D. Sciti, Efficacy of a ZrB2-SiC matrix in protecting C fibres   60 Accepted Manuscript  continuous carbon fibre reinforced ZrB2-SiC composite, Corros. Sci. 123   [26] Y. Miaoa, F. Zhanga, Z. Yangb, D. Jiab, Incorporation of BN-coated   [25] A. Vinci, L. Zoli, E. Landi, D. 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  "_id": 11,
  "PDF": "Ablation Property of ZrB2-SiC Composite Sharp Leading Edges with Varying Radiuses of Curvature under Oxy-Acetylene Torch.pdf",
  "Text": "['Key Engineering Materials Vols. 512-515 (2012) pp 710-714 © (2012) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.512-515.710  Online: 2012-06-04  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, www.ttp.net. (ID: 128.210.126.199, Purdue University Libraries, West Lafayette, USA-07/07/15,12:22:40)  Ablation Property of ZrB2-SiC Composite Sharp Leading Edges with Varying Radiuses of Curvature under Oxy-acetylene Torch Rujie He a, Xinghong Zhang and Ping Hu National Key Laboratory of Science and Technology on Advanced Composites in Special Environments, Harbin Institute of Technology, Harbin 150001, P. R. China a herujie2003jci@163.com Keywords: ZrB2-SiC composite; Sharp leading edges; Radius of curvature; Ablation. Abstract. Ablation behavior of ZrB2-SiC sharp leading edges with five different curvature radiuses was investigated using an oxy-acetylene torch. During the test, the curvature radiuses were 0.15 mm, 0.5 mm, 1.0 mm, 1.5 mm, and 2.0 mm, respectively. Under the same ablation condition, the smaller was the radius, the severer ablation underwent. The sharp leading edge with a curvature radius of 0.15 mm had the highest surface temperature and maximum surface temperature rising rate, exceeded 2100 ºC in less than 30 s. However, the surface of sharp leading edge with a curvature radius of 2.0 mm achieved only 1900 ºC in more than 60 s. After 5 min ablation, the mass and linear ablation rates were measured. All the five sharp leading edges evolved to nearly a same radius after ablation. The microstructure of the oxidation layers was also investigated. A ZrO2-SiO2 layer generated from oxidation of ZrB2-SiC acts as a thermal barrier and reduces the diffusion of oxygen.  Introduction Ultra high temperature ceramics (UHTCs), including some refractory metal diborides, such as zirconium diboride (ZrB2), have been historically studied and developed since 1960s and early 1970s [1-3]. These materials are of great interest for their unique properties combined with ultra-high melting temperature (>3000 ºC), high electrical and thermal conductivities, as well as excellent chemical inertness, which make them the most promising candidates for ultra-high temperature applications, such as sharp leading edges and nosecones, and thermal protection systems for reusable atmosphere re-entry vehicles, hypersonic flight, and rocket propulsion systems [4-12]. Recent interest of these ceramics has focused on densification, mechanical properties, thermal shock and oxidation behavior, and ablation resistance property.  As a critical enhancement, oxidation and ablation behavior of ZrB2-based ceramics has been investigated widely by many researchers. SiC is typically added to monolithic ZrB2 ceramics to improve the strength, fracture toughness and oxidation resistance. For temperature above 1200 ºC, the addition of SiC provides more efficient oxidation resistance by encouraging the formation of borosilicate glass on the exposed surfaces, providing a much greater oxidation protection [13-16].  Furthermore, in contrast to traditional blunt capsules or shuttle-type vehicles, characterized by poor gliding capabilities, the future of UHTCs opens new prospects for development of space planes with slender fuselage noses and sharp wing leading edges. Vehicles with sharp leading edges and/or sharp noses have lower drag and high lift to drag ratios (L/D) than blunt-nosed vehicles, but they also have to endure higher surface temperatures [17-20]. As for the application of sharp leading edges and/or sharp noses, low radius leading edges are subjects to much areothermal heating than blunt edges, such as those on the space shuttle orbiter, and these edges will reach temperatures that may exceed 2000 ºC during re-entry conditions. Sharp leading edges have not been used for reusable vehicles in the past because materials able to repeatedly withstand the higher entry and abort temperature were not available. Recent developments in the use of UHTCs have given renewed hope to the eventual realization of sharp leading edge vehicles. However, to our best knowledge, virtually nothing has been published in the open literature discussing the ablation behavior of UHTC sharp leading edges with different radiuses. In this study, sharp leading edges prepared by ZrB2-20SiC ultra high temperature ceramics with different curvature \\x0c', 'Key Engineering Materials Vols. 512-515  711   radiuses were tested using an oxyacetylene torch flame under the same condition for comparison. The purpose of this study is to investigate the ablation behavior and microstructure evolution of the ZrB2-SiC composite with different radiuses and the ablation mechanism is also discussed.  Table 1. Experimental conditions for oxyacetylene ablation. O2 gas pressure (kPa) 0.45 O2 gas flux (L/s) 0.6 C2H2 gas pressure (kPa) 0.1 C2H2 gas flux (L/s) 0.4 Diameter of nozzle (mm) 2 Distance from sample surface to nozzle (mm) 20 Experimental procedure Commercial ZrB2 (2 µm; purity>99.5%; Northwest Institute for Non-ferrous Metal Research, China) and SiC (0.5 µm; purity>99.5%; Weifang Kaihua Micro-powder Co. Ltd., China) powders were used in this study. The powder of ZrB2-20vol%SiC was ball-milled in ethanol for 8 h using zirconia media and dried in a rotating evaporator. Milled powders were uniaxially hot pressed in a boron nitride coated graphite die at 2000 ºC for 1 h under vacuum (0.5 mbar) and 30 MPa of applied pressure. The heating schedule has been described in detail elsewhere [13-16]. Ablation behavior was evaluated by means of the oxyacetylene torch. Table 1 lists the specific experimental condition. During the ablation tests, ZrB2-SiC sharp leading edge models with different curvature radiuses (Fig. 1) were cut from the biller. The curvature radiuses are 0.15 mm, 0.5 mm, 1.0 mm, 1.5 mm, and 2.0 mm, respectively. The corresponding samples were designated as R015, R050, R100, R150, and R200 for simplicity. Ablation time was chosen as 600 s. Surface temperature of the ceramic sharp leading edges was determined using a two-color Raytek pyrometer (RAYMR1SCSF, USA). Crystalline phases and microstructure of the specimen after ablation were indentified by X-ray diffraction (XRD; Rigaku, Dmax-rb, Cu Kα=1.5418 Å) and scanning electron microscopy (SEM; FEI Sirion, Holland), respectively. The mass and linear ablation rates of the samples were also obtained. Results and discussion The bulk density of the hot-pressed ZrB2-SiC composite was 5.41 g/cm3, which corresponds to a relative density higher than 98%. Fig. 2 is the microstructure of the polished surface of ZrB2-SiC composite. The SiC particulates, which represent the majority of the dark features, are homogeneously dispersed in the ZrB2 matrix and no agglomeration was detected.  Fig. 1. Schematic diagram and macrographs of ZrB2-SiC sharp leading edges with different curvature radiuses before and after ablation (a, b, c, f, and e, macrographs before ablation for R015, R050, R100, R150, and R200, respectively; f, g, h, i, and j, macrographs after ablation for R015, R050, R100, R150, and R200, respectively).  Fig. 2. Microstructure of the polished surface of the ZrB2-SiC composite  \\x0c', '712  High-Performance Ceramics VII   As shown in Table 2, the sharp leading edges with smaller radiuses exhibited severer ablation when ablated by the oxyacetylene torch. When the radiuses are 0.15mm and 0.50mm, the mechanical erosion is much severer than the weight gain and volume expansion caused by the oxidation of the material, so the mass and linear ablation rates are positive values. While the radiuses are 1.00mm, 1.50mm and 2.00mm, the mechanical erosion can not move away the materials due to the bigger radiuses, and the oxidation of ZrB2-SiC become more significant leading to the weight gain and volume expansion, resulting in the mass and linear ablation rates are negative values.  Table 2. The ablation properties of the ZrB2-SiC samples with different radiuses Label Radius (mm) Ablation time (s) Mass ablation rate (g/s) Linear ablation rate (mm/min) Surface temperature (ºC) R015 0.15 600 +3.3×10-5 * +1.8×10-3 2100 R050 0.5 600 +1.7×10-5 +4.7×10-4 2040 R100 1.0 600 -3.0×10-5 -2.7×10-4 2000 R150 1.5 600 -1.34×10-4 -4.3×10-4 1930 R200 2.0 600 -1.70×10-3 -8.3×10-4 1900 * + refers to weigh loss and linear shrinkage; refers to weight gain and linear expansion.   Fig. 3. Surface temperature curves vs. time (s) during ablation of ZrB2-SiC sharp leading edges with different curvature radiuses.  Fig. 3 plots the temperature curves on the surface of the five ZrB2-SiC sharp leading edges varied with the ablation time increased up to 300 s. Sharp increments in the surface temperature could clearly be found for all samples with different radiuses at the beginning of ablation. Hereafter, the temperature reached a steady state with the maximum points as ~2100 ºC for R015, ~2040 ºC for R050, ~2000 ºC for R100, ~1930 ºC for R150, and ~1900 ºC for R200, respectively. Furthermore, the surface temperature of R015 reached up to the maximum in less than 30 s, much quicker than that of other samples (~32 s for R050, ~43 s for R100, ~50 s for R150, and ~60 s for R200, respectively). Such difference in the surface temperature was probably ascribed to the different radiuses of these samples.  Besides, it was observed from Fig.3 that the surface temperature of all samples trended to nearly a same temperature (around ~1950 ºC) at last, indicating that the samples with different radiuses were ablated to nearly a same radius (R≈1.5mm) and finally resulting in nearly a same surface temperature (T≈1950 ºC) after a long time ablation. Fig. 4 shows the microstructure of the cross section of the samples with different radiuses after 600s oxyacetylene torch testing. It can be clearly found that all the ZrB2-SiC sharp leading edges with different radiuses convert to nearly a same radius, about 1.5 mm. The oxidation layer of the ZrB2-SiC \\x0c', 'Key Engineering Materials Vols. 512-515  713   ceramic is consisted of four distinct layers. The outmost layer was rather compact except for a few large pores, and previous research results indicated that this layer was mainly composed of Si, O, and a little Zr. Some zirconia coalesced together by recrystallization under the present experimental condition. SiO2 glass formed by the oxidation of SiC was embedded within the compact ZrO2. The formation of this layer effectively inhibited inward transport of oxygen. Passive oxidation protection was provided. Underlying this outmost layer was a porous crystalline ZrO2 rich layer, which was a transitional region between the SiC-depleted layer and the outmost oxide layer. Underneath, a partially depleted SiC layer was formed, which was caused by the active oxidation of SiC. The forth oxide layer is the unaltered ZrB2-SiC material (see Fig. 5). It was observed that the thickness of the glass SiO2 layer was, for the difference of the surface temperature and strong mechanical erosion of these samples, causing the gasification, evaporation and move-away of glass SiO2. Analysis about the microstructure indicated that the surface of the sharp leading edges were covered with a thick layer of SiO2 and ZrO2, although the glass SiO2 may evaporated due to the high temperature. It was well known that oxides produced from the oxidation of ZrB2 and SiC were able to fill and close the microdefects and then resisted the oxygen flow, which acted as a thermal barrier and reduces the diffusion of oxygen.    Fig. 4. Scanning electron micrographs (SEM) cross-section of the samples with different radiuses after 600s ablation: (a) R015; (b) R050; (c) R100; (d) R150; (e) R200.   Fig. 5. Microstructure of the oxidation layers after ablation: (1) glass SiO2 layer; (2) rich ZrO2 layer; (3) SiC depletion layer and (4) the unaltered ZrB2-SiC material. Conclusions By means of the oxyacetylene torch, ablation behavior of ZrB2-SiC sharp leading edges with different radiuses was investigated. Results of surface temperature showed that sample with smaller radius reached higher surface temperature and quicker temperature razing rates. Particularly, all the sharp leading edges with different radiuses convert to nearly a same radius. The oxidation layers were consisting of four layers: (1) a porous glass SiO2 layer; (2) a crystalline ZrO2 rich outer layer; (3) SiC depletion layer, and (4) unaltered ZrB2-SiC material. A ZrO2-SiO2 layer generated from oxidation of ZrB2-SiC acts as a thermal barrier and reduces the diffusion of oxygen. \\x0c', '714  High-Performance Ceramics VII   Acknowledge This research was supported by the National Science Foundation (51072042 and 50972029) of China, project (HIT. KLOF. 2009026) supported by the Key Laboratory Opening Funding of National Key Laboratory on Advanced Composites in Special Environment. References [1] A.K. Kuriakose, J.L. Margrav, The Oxidation Kinetics of Zirconium Diboride and Zirconium Carbide at High Temperatures, J. 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Soc. 22 (2002) 279-288.  http://dx.doi.org/10.1016/S0955-2219(01)00284-9  [5] F. Monteverde, The Thermal Stability in Air of Hot-pressed Diboride Matrix Composites for Uses at  Ultra-high-temperatures, Corros. Sci. 47 (2005)2020-(2033).  http://dx.doi.org/10.1016/j.corsci.2004.09.019  [6] S. Norasetthekul, P.T. Eubank, W.L. Bradley, B. Bozakurt, B. Stucker, Use of Zirconium Diboride Copper as an Electrode in Plasma Applications, J. Mater. Sci. 34 (1999) 1261-1270.  http://dx.doi.org/10.1023/A:1004529527162  [7] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, High Strength Zirconium Diboride Based Ceramics, J. Am. Ceram. Soc. 87 (2004) 1170-1172.  http://dx.doi.org/10.1111/j.1551-2916.2004.01170.x  [8] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J. A. Zaykoski, Refractory Diborides of Zirconium and  Hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364.  http://dx.doi.org/10.1111/j.1551-2916.2007.01583.x  [9] F. Monteverde, R. Savino, Stability of Ultra-high-temperature ZrB2-SiC Ceramics under Simulated  Atmospheric Re-entry Conditions, J. Eur. Ceram. Soc. 27 (2007) 4797-4805.  http://dx.doi.org/10.1016/j.jeurceramsoc.2007.02.201  [11] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of Ultra-High  Temperature Ceramics for Aeropropulsion Use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767.  http://dx.doi.org/10.1016/S0955-2219(02)00140-1  [13] P. Hu, G.L. Wang, Z. Wang, Oxidation Mechanism and Resistance of ZrB2-SiC Composites, Corro. Sci.  51 (2009) 2724-2732.  http://dx.doi.org/10.1016/j.corsci.2009.07.005  [14] X.H. Zhang, P. Hu, J.C. Han, et al., Ablation behavior of ZrB2-SiC ultra high temperature ceramics  under simulated atmosphere Re-entry conditions, Comp. Sci. Techno. 68 (2008) 1718-1726.  http://dx.doi.org/10.1016/j.compscitech.2008.02.009  [15] J.C. Han, P. Hu. X.H. Zhang, S.H. Meng, W.B. Han, Oxidation Resistant ZrB2-SiC Composites at 2200  ºC, Comp. Sci. Techno. 68 (2008) 799-806.  http://dx.doi.org/10.1016/j.compscitech.2007.08.017  [16] J.C. Han, P. Hu. X.H. Zhang, S.H. Meng, Oxidation Behavior of Zirconium Diboride-Silicon Carbide at        \\x0c', '1800 ºC, Scripta Mater. 57 (2007) 825-828.  http://dx.doi.org/10.1016/j.scriptamat.2007.07.009  [17] R. Savino, F.M. De Stefano, D. Paterna, M. Serpico, Aerothermodynamic Study of UHTC-based  Thermal Protection Systems, Aero Sci. Tech. 9 (2005) 151-160.  http://dx.doi.org/10.1016/j.ast.2004.12.003   \\x0c']"
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  "_id": 12,
  "PDF": "Ablation resistance and mechanism of SiC–LaB6 and SiC–LaB6–ZrB2 ceramics under plasma flame.pdf",
  "Text": "['Ceramics International 46 (2020) 16249-16256  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Ablation resistance and mechanism of SiC-LaB6 and SiC-LaB6-ZrB2 ceramics under plasma flame  Hanzhou Liua, Xin Yanga,b,∗, Cunqian Fanga, Anhong Shia, Lei Chena, Qizhong Huanga  T  a State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China b Science and Technology of Advanced Functional Composite Laboratory, Aerospace Research Institute of Materials & Processing Technology, Beijing, 100076, PR China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ablation Spark plasma sintering Silicon carbide Lanthanum hexaboride Zirconium boride  1.  Introduction  In this study, silicon carbide-lanthanum hexaboride (SiC-LaB6) and silicon carbide-lanthanum hexaboride-zirconium boride (SiC-LaB6-ZrB2) ceramics were fabricated by spark plasma sintering at 1900 °C, and their ablation resistance was tested under plasma flames over 2300 °C. The results indicate that the SiC-LaB6-ZrB2 ceramic exhibits better ablation resistance than the SiC-LaB6 ceramic. After ablation under the plasma flame for 60 s, the mass and linear ablation rates of the SiC-LaB6 ceramic were 15.83 μg/s and 1.08 μm/s, respectively, while those of SiC-LaB6-ZrB2 were -8.42 μg/s and -0.27 μm/s. With the addition of ZrB2, SiC-LaB6-ZrB2 ceramic attained a high density and fewer inner oxygen diffusion channels. Moreover, the ZrO2-La2O3-SiO2 oxide scale with good self-healing ability and excellent stability was formed in the ablation centre, which can retard the further oxidation during ablation.  Silicon carbide (SiC) is a widely used ultrahigh temperature ceramic (UHTC) that has attracted attention of many researchers for its high melting point (> 2800 K), low thermal expansivity (4.5 × 10−6/°C) and excellent oxidation resistance [1,2]. Under the high temperature (1700 °C) oxidizing environments, SiC can form self-healing silica glass with low oxygen permeability, thereby effectively preventing the further diffusion of oxygen [3-6]. However, once the temperature exceeds 1700 °C, the protection of silica glass will be damaged owing to volatilization of the formed silicon dioxide (SiO) [7,8]. Therefore, the applications of monolithic SiC ceramic in the latest generation of hypersonic vehicle and re-entry aircraft are restricted owing to its unsatisfactory oxidation and ablation properties at ultra-high temperature. To promote the ablation resistance of SiC, numerous attempts by adding additives have been reported by researchers. Recently, Lanthanum hexaboride (LaB6) is regarded as a promising additive for its advantages of low thermal expansion coefficient, high hardness, and good chemical stability [9,10]. It is reported that adding LaB6 can ef[11-14]. Acfectively improve the oxidation resistance of UHTCs cording to the results of the previous study [15], the generation of La2Si2O7 molten phase in La2O3 modified C/C-SiC composites during ablation can enhance the ablative properties of the composites.  However, in severe ablative environments with high temperatures and violent scouring, the loss of glass film or molten phases with low viscosity will induce material failure. Therefore, in particular for UHTCs, the stable oxide barrier formed during ablation is vital for promoting ablation resistance. Zirconium boride (ZrB2) is an ideal reinforcing phase owing to its high melting temperature, moderate density, low thermal expansion coefficient and high thermal conductivity [16]. Moreover, the oxidation product ZrO2 can maintain the integrity of the ablated surface and retard the further pervasion of the oxygen on account of its low vapor pressure at high temperatures [17,18]. For instance, the B2O3 provided by ZrB2 exhibits a high wettability and considerable surface tension, which can seal the cracks in the ceramic, thereby improving the oxidation resistance at medium temperature (800-1200 °C) [19,20]. In addition, ZrB2 combines well with SiC and exhibits an outstanding oxidation resistance [21-25]. In this work, SiC-LaB6 and SiC-LaB6-ZrB2 ceramics were prepared by spark plasma sintering process (SPS) at 1900 °C. To obtain a highdensity SiC-LaB6-ZrB2 ceramic, ZrB2 phase was generated through the in situ reaction during the sintering process. The microstructure and density of the sintered ceramics were analyzed in detail, and the ablation behaviour and mechanisms of both ceramics were investigated under plasma flame (approximately 2300 °C).  ∗ Corresponding author. State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China. E-mail address: yangxincsu@csu.edu.cn (X. Yang).  https://doi.org/10.1016/j.ceramint.2020.03.181 Received 9 December 2019; Received in revised form 21 February 2020; Accepted 18 March 2020  Available online 20 March 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'H. Liu, et al.  Table 1  Density of SL and SLZ sintered at 1900 °C.  Composition (vol%)  Theoretical density (g/cm3)  Bulk density (g/cm3)  Relative density (%)  SiC-40%LaB6 SiC-25%LaB6-15% ZrB2  3.96 4.41  3.18 3.91  80.4 88.6  2. Experimental  2.1. Preparation  Commercially available ZrC powder (average particle size of 2 μm, purity of > 99.5%), LaB6 powder (average particle size of 2 μm, purity size of 1 μm, purity of > 99.5%) and SiC powder (average particle SiC-LaB6 of > 99.5%) were used as raw powders to prepare and SiC-LaB6-ZrB2 ceramics. The following compositions were produced: SiC +40 vol%LaB6 and SiC +25 vol%LaB6 + 15 vol%ZrC. These powder mixtures were ball milled with ZrO2 balls at 400 rpm for 8 h in absolute ethylalcohol with ZrO2 milling media. After drying by rotary evaporation, the two prepared powder mixtures were loaded into graphite die (40 mm in diameter) and hot pressed via SPS equipment (Model HPS-200, Kingtier New Alloy Material Co. Ltd, Anhui, China) in a vacuum environment at 40 MPa; the temperature was increased at a rate of approximately 100 °C/min, and the target temperature 1900 °C was maintained for 5 min. Finally, the obtained samples were cooled down to room temperature at a rate of 40 °C/min in the furnace. The ceramic samples manufactured in the volume proportions SiC-40%LaB6 and SiC-25%LaB6-15%ZrC are denoted as ‘SL’ and ‘SLZ’, respectively.  2.2. Ablation test  The ablation test was conducted with Multiplaz 3500 plasma ablation equipment. The surface of the samples was held perpendicular to the plasma gun at a distance of 10 mm and the inner diameter of the nozzle tip was 2 mm. The flame temperature was measured with an optical pyrometer (2300 °C). The specific mass ablation rate (mg/s) is calculated by measuring the mass before and after ablation with the following equation:  R  m  =  m  2  m  1  −  t  (1)  where the Rm is the mass ablation rate (mg/s), m1 the mass (mg) of the specimen before ablation, m2 the mass after ablation, and t is the ablation time (s). The specific linear ablation rate (mm/s) is calculated by measuring the thickness before and after ablation with the following equation:  R  l  =  l  2  l  1  −  t  Ceramics International 46 (2020) 16249-16256  (2)  where Rl is the linear ablation rate (mm/s), l1 and l2 are the thicknesses (mm) of the specimen before and after ablation, respectively, and t is the ablation time (s).  2.3. Characterization  The bulk density of the samples was measured by using the Archimedes method, and the theoretical densities of the samples were calculated by applying the rule of mixture. Finally, the relative density was determined by dividing the bulk density by the theoretical density. The crystalline structure and morphology of the samples before and by X-ray after ablation were examined diffraction (Rigaku Dmax/ 2550VB+18 KW) and scanning electron microscope (NOVA NanoSEM230, Czech Republic) equipped with an energy dispersive spectroscope (EDS, EDAX Inc). The selected area electron diffraction (SAED) pattern was obtained by transmission electron microscopy (TEM, FEI, Tecnai G2 F20).  3. Results and discussion  3.1. Microstructure and density of SiC-LaB6 and SiC-LaB6-ZrB2 ceramics  The densities of SL and SLZ are listed in Table 1. The theoretical densities were calculated with 3.22 g/cm3 for SiC, 4.61 g/cm3 for LaB6, and 6.57 g/cm3 for ZrC along with the rule of mixture, and the bulk densities were measured by the Archimedes method. According to Table 1, the relative density of SLZ is approximately 10% higher than that of SL. Thus, the density of SiC-LaB6 ceramic has been remarkably enhanced by the addition of ZrB2. Fig. 1 presents the SEM images of SL and SLZ. As shown in Fig. 1(a), the microstructure of SL is compact without evident defects, and the the ceramic particles are in the range of 2-5 μm. Thus, dimensions of the LaB6 and SiC crystal grains grow up during the SPS process. As shown in Fig. 2(b), the ceramic particles in SLZ(1-2 μm), are evidently smaller than those of SL. Owing to the uniform and fine dimension, the combination of the ceramic particles in SLZ has been remarkably enhanced. Therefore, fewer voids has been formed inside SLZ, and the ceramic exhibits a relatively high density. The phase compositions of SL and SLZ are presented in Fig. 2. As shown in Fig. 2(a), the dominant phases of SL are LaB6 and SiC, which are the same as those of original powders without any new phases or impurities. By contrast, SLZ is composed of LaB6, SiC, ZrB2, and La2C3. Thus, after the SPS process, ZrC has converted completely into ZrB2 and La2C3 has been generated owing to the chemical reaction between LaB6 and ZrC. To identify the phase compositions precisely, TEM analyses were  Fig. 1. SEM images of polished ceramic microstructures: (a) SL and (b) SLZ.  16250  \\x0c', 'H. Liu, et al.  Ceramics International 46 (2020) 16249-16256  Fig. 2. Phase compositions of SL and SLZ: (a) XRD patterns; (b) TEM of La2C3; (c)SAED pattern of La2C3.  Table 2  Comparison of measured and calculated data (d and ϕ) for La2C3.  Ref.  -  PDF:820622  Measured Calculated  d(221) (nm)  d(122) (nm)  d(301) (nm)  ϕ < (221), (122) > (°)  ϕ < (122), (301) > (°)  0.2929 0.2939  0.2943 0.2939  0.2797 0.2788  63.45 63.61  58.30 58.19  La2C3 can be obtained: the corresponding results are listed in Table 2. According to Table 2, the measured and calculated values exhibit no remarkable differences. Hence, the SLZ contains La2C3. The equations of d and ϕ for the cubic system are presented as below, respectively:  d  =  a  2  h  2  k  +  2  l  +  cos  ϕ  =  h h  2  h  2 1  k  +  1 2 1  +  k k  1  2  +  l  l  1 2  h  2 2  k  2 2  +  l  2 2  +  +  +  l  2 1  (3)  (4)  The previously presented analysis confirms the generation of ZrB2 and La2C3. It can be concluded that the formation of refined grains in SLZ is probably attributed to the in situ reaction between LaB6 and ZrC, which has substantially reduced the void number, and has promoted the ceramic density.  3.2. Ablation morphology of SiC-LaB6 and SiC-LaB6-ZrB2 ceramics  The optical photographs of the ablated surfaces of SL and SLZ are shown in Fig. 3. According to Fig. 3(a), the SL surface is compact without visible cracks and pits. Thus, SL has a good ablation resistance under the violent scouring force of the plasma flame. A white circle (approximately 5 mm in diameter) has occurred on the ablated central area of SL, which is the severely ablated region that has faced the plasma flame directly. Except for the central area, the rest of the ablated surface is pale-white owing to the slight oxidation of the ceramic. As shown in Fig. 3(b), the ablated surface appearance of SLZ is similar to that of SL. However, the white circle in the central area of SLZ is less pronounced. Thus, the addition of ZrB2 has a positive effect on the ablation resistance: it protects the ceramic from further oxidation.  Fig. 3. Optical photographs of the ablated ceramic surfaces: (a) SL and (b) SLZ.  conducted by TEM. As shown in Fig. 2(b), the maximal diameter of the ceramic particle is approximately 500 nm, which is consistent with the broad diffraction peak of La2C3. With the SAED pattern (Fig. 2(c)), crystal parameters of La2C3 (PDF:82-0622), and the equations of the cubic system (Eqs. (3) and (4)), the measured and calculated data of  16251  \\x0c', 'H. Liu, et al.  Ceramics International 46 (2020) 16249-16256  Fig. 4. Center ablated surface morphology of SL: (a) SEM of center point and outer ablation region; (b) magnified SEM of center point; (c) BSED of molten phase; (d) EDS result of part 1; (e) EDS result of part 2.  Fig. 4 shows the surface morphology of SL after ablation. According to Fig. 4(a), the ablated SL sample has survived the plasma ablation test; only few cracks and pits are visible on the ablated surface. Furthermore, the ablated surface is covered by numerous molten phases, and the oxide scale consisting of molten phases is porous and trends to bulge. The outer ablation region exhibits fewer pores than the center point. Thus, more oxide gases have been released from the central area owing to severe ablation. The high-magnification SEM image (Fig. 4(b)) presents smooth molten phases in the center point, which reveals their good flowability. During the ablation, considerable amounts of cracks and pits have been healed up by molten phases. Thus, the oxide scale on the ablated surface possesses self-healing ability. However, owing to the  high pressure and strong heat flow provided by the plasma flame, more pores have formed in the center point because a large amount of oxide gases was released under the ablated surface. In addition, the evaporation rate of the molten phases in the central area was accelerated, and the insufficient number of molten phases could not seal the pores. According to BSED image (Fig. 4(c)), the molten phases on the ablated surface contains square and acicular grains. This interlaced microstructure of molten phases helps preventing excess evaporation by reducing the flowability. The EDS results (Fig. 4(d and e)) indicate that the molten phase is composed of La-Si-O compounds. The Au occurrence originates from the gold spray treatment before the SEM investigation.  16252  \\x0c', 'H. Liu, et al.  Ceramics International 46 (2020) 16249-16256  Fig. 5. Center ablated surface morphology of SLZ: (a) SEM of center point and outer ablation region; (b) magnified SEM of center point; (c) BSED of oxide scale; (d) EDS analysis of part 1; (e) EDS result of part 2; (f) EDS result of part 3.  Fig. 5 shows the surface morphology of SLZ after ablation. In Fig. 5(a), the ablated central area of SLZ is flat occur, and only few cracks and pits can be observed. In addition, only few molten phases and pores occur on the ablated SLZ surface, which distinguishes it from the ablated SL surface (Fig. 4(a)). Fig. 5(b) shows that the ablated SLZ surface is covered by an oxide scale comprised of substantial particles and few molten phases. In addition, owing to the tight combination of oxide products, the oxygen diffusion tunnels on the ablation surface are reduced, which have prevented the further oxidation of the SLZ ceramic. The compact and dense ablated surface reveals that SLZ has a remarkable ablation resistance. According to the BSED image (Fig. 5(c)) and EDS analysis (Fig. 5(d)), the skeleton of the oxide scale (dark-grey phases; NO. 2 in Fig. 5(c)) is composed of O, Si, Zr and La. The light grey phases contain a low Zr amount (NO. 1 and 3 in Fig. 5(c)), which occurs mostly in the molten state. Therefore, the viscosity of the oxide scale has been promoted with the addition of a Zr-based UHTC phase, which has stabilized the ablated surface under the violent brunt of the plasma flame. Furthermore, the molten phases play an important role in healing up cracks and pits on the ablated surface. XRD analyses were conducted to investigate the oxide compositions of SZ, and SLZ more thoroughly; the results are shown in Fig. 6. According to Fig. 6(a), the dominate oxidation products in SL are La2Si2O7 and La2O3. The low melting point (approximately 1750 °C) causes the La2Si2O7 molten phases to have a low viscosity with self-healing ability. Together with the highly viscous La2O3 phases (Fig. 4(c)), the evaporation of La2Si2O7 molten phases has been restrained. Therefore,  16253  \\x0c', 'H. Liu, et al.  Ceramics International 46 (2020) 16249-16256  Fig. 6. XRD patterns of central ablated surfaces: (a) SL and (b) SLZ.  Table 3  Ablation rates of SL and SLZ for 60 s.  Composition (vol%)  Mass ablation rate (μg/s)  Ablated-centre linear ablation rate (μm/s)  SiC-40%LaB6 SiC-25%LaB6-15%ZrC  15.83 −8.42  1.08 −0.27  Table 4  Mass ablation rates on the specific surface area reported in other literature.  Sample  Ablation time (s)  mass ablation rate on the specific surface area (mg·cm−2 ·s−1)  SLZ C/C-SiC [27] C/C-HfC [28] C/C-ZrB2-ZrC-SiC [29] ZrC-SiC coated C/C-ZrC-SiC [30]  60 30 240 180 120  −0.043 5.74 0.55 5.09 22.93  La2Si2O7 molten phase has survived on the ablated surface, and the ablation resistance has been improved accordingly. As shown in Fig. 6(b), m-ZrO2, t-ZrO2, La2O3, and SiO2 were detected in SLZ. The tZrO2 has formed after rapid cooling of the molten particles from ultrahigh temperature to room temperature. The peak intensity of La2O3 is weak. However, the EDS results of the ablated surface (Fig. 5(d)-(f)) indicates that the La content is not low. According to the relevant report [26], it is possible that a part of La2O3 occurred in the solid solution in the formed ZrO2, which has contributed to the formation of t-ZrO2. Compared with that of the single m-ZrO2 phase, the t-ZrO2 and m-ZrO2 polyphase has a high toughness because the energy that promotes crack growth is absorbed by the transition from t-ZrO2 to m-ZrO2 (martensite transformation). Therefore, the oxide scale on the ablated SLZ surface has maintained a compact structure owing to the t-ZrO2 phase even after the rapid cooling process. Moreover, the viscosity of the oxide scale has been adjusted by the formed SiO2 molten phase. Thus, defects have been effectively sealed on the ablated surface.  3.3. Ablation rates  The ablation rates of SL and SLZ for 60 s are listed in Table 3. The mass and linear ablation rates of SL are both positive, which is caused by the rapid evaporation of B2O3 glass and the partial loss of La2Si2O7 in the central ablated area. Despite the loss of B2O3 and SiO2 glass, the ablation rates of SLZ are both negative, which demonstrates the extraordinary stability of ZrO2 and La2O3 under the plasma flame. Both  mass and linear ablation rates of SLZ are evidently lower than those of SL. The enhanced mass and linear ablation rates of SLZ can be interpreted as following: On the one hand, the addition of ZrB2 improves the density of SLZ, which reduces the oxygen inner diffusion channels inside SLZ. On the other hand, with the addition of ZrB2, ZrO2-SiO2-La2O3 molten phases with high viscosity and stability were formed on the ablated surface of SLZ. Due to the high melting point (about 2700 °C) of ZrO2, the viscosity and anti-scouring behavior of the ZrO2-SiO2-La2O3 molten phases were improved with incorporation of ZrO2 skeleton. Therefore, the scouring of SiO2 glass was retarded and the ZrO2-SiO2-La2O3 molten phases remained stable during the ablation, which effectively enhances the ablation resistance of SLZ. Table 4 displays the reported mass ablation rates on the specific surface area of C/C-UHTC composites under the similar plasma flame condition. Considering the effective ablated area in the ablation center (approximately 5 mm in diameter), the calculated mass ablation rate of SLZ on the specific surface area is -0.043 mg cm−2 ·s−1. According to Table 4, the mass ablation rate of SLZ on the specific surface area is negative, while those of C/C-UHTC composites are positive, indicating the superior ablation property of SLZ ceramic than C/C-UHTC composites. SiC ceramics exhibit excellent oxidation resistance under 1600 °C [1-6]. However, with the increasing ablation temperature (> 2200 °C), the active oxidation of SiC can generate large amounts of oxide gases (SiO and CO), damaging the self-healing ability of the silica glass [7,8]; meanwhile, the formed SiO2 glass with low viscosity will vaporize and is easily scoured away under the high temperature flame. Therefore, the ablation resistance of SiC ceramics decreases rapidly above 2200 °C. The negative mass and linear ablation rates of SLZ in this work demonstrate that adding UHTCs can improve the ablation resistance of SiC ceramics at high temperature.  3.4. Ablation mechanism  Based on the previous analysis, it can be confirmed that the oxidation products and surface morphology of the ablated two samples are different. The ablation mechanisms of SL and SLZ may be as follows: the SiC-40%LaB6 ceramic reacts with oxygen in the beginning of the ablation procedure, and La2O3, SiO2 and B2O3 phases are generated on the surface (Eqs. (5) and (6)). Owing to the good wettability and low viscosity of B2O3 glass, the defects on the ceramic surface are healed up and the inner diffusion of oxygen is delayed. Therefore, the oxidation resistance of SL at medium temperatures (800-1200 °C) is improved. However, when the ablation temperature exceeds 1200 °C, the B2O3 glass evaporates (Eq. (7)). In addition, La2Si2O7 is generated and fuses into molten scale at about 1700 °C (Eqs. (8) and (9)). The La2Si2O7 molten phase has a similar self-healing effect as B2O3 glass. Moreover,  16254  \\x0c', 'H. Liu, et al.  Ceramics International 46 (2020) 16249-16256  the viscosity of the molten La2Si2O7 increases with the doping of La2O3, which partially impedes the scouring of the oxide scale. Nevertheless, after exposure to the strong plasma flame for a long time, the La2Si2O7-La2O3 molten phases are consumed by the constant evaporation and scouring. Due to the inferior stability of La2Si2O7-La2O3 molten phase at high temperatures (≥2300 °C), the ablation resistance of the SL ceramic is limited to a certain extent.  LaB6(s) + O2(g) → La2O3(s) + B2O3(l)  SiC(s) + O2(g) → SiO2(s) + CO(g)  B2O3(l) → B2O3(g)  La2O3(s) + SiO2(s) → La2Si2O7(s)  La2Si2O7(s) → La2Si2O7(l)  (5)  (6)  (7)  (8)  (9)  The original ablative behaviour of SLZ is similar to that of SL, and ZrO2, La2O3, SiO2, and B2O3 are initially formed (Eqs. (10)-(12)) on the ablated surface, which withstands the violent scouring from the plasma flame and maintains the integrity of the oxide scale during the entire ablation process. After the rapid loss of B2O3 glass, SiO2 transforms into liquid glass ((Eq. (13), approximately 1750 °C), which promotes the self-healing ability of the oxide scale continuously. Usually, pure SiO2 glass is easily scoured away from the surface as B2O3 glass. However, owing to the moderate melting point and viscosity of ZrO2-SiO2-La2O3 molten phases, the scouring of SiO2 glass is retarded and the defects in the ZrO2-La2O3 skeleton can be effectively sealed. In the ablation ZrO2-La2O3-SiO2 process, the formed oxide scale with good selfhealing ability and high stability can act as a thermal barrier that protects the substrate from further ablation. Because of the compact inner matrix with the high density, the inner oxygen diffusion tunnels inside the SLZ are reduced, which retards effectively the further oxidation of the ceramic. Therefore, with the addition of the Zr-based UHTC phase, the ablation resistance performance of SLZ is significantly improved.  ZrB2(s) + O2(g) → ZrO2(s) + B2O3(l)  La2C3(s) + O2(g) → La2O3(s) + CO(g)  SiC(s) + O2(g) → SiO2(s) + CO(g)  SiO2(s) → SiO2(l)  4. Conclusion  (10)  (11)  (12)  (13)  In this study, SiC-LaB6 and SiC-LaB6-ZrB2 ceramics were fabricated by SPS, and the effects of ZrB2 on the ablation resistance of the SiC-LaB6-ZrB2 ceramic were investigated. With the addition of ZrB2, the density of SiC-LaB6-ZrB2 (88.6%) was remarkably enhanced compared with that of SiC-LaB6 (80.4%). This can be attributes to the grain refinement during the in situ reaction between LaB6 and ZrC. Furthermore, a ZrB2 phase was generated in the reactive sintering process. After 60 s plasma flame ablation, the mass and linear ablation rates of the SiC-LaB6 ceramic were 15.83 μg/s and 1.08 μm/s, respecμg/s tively; those of the SiC-LaB6-ZrB2 ceramic were -8.42 and μm/s, -0.27 respectively. During ablation, La2Si2O7-La2O3 molten phases with self-healing ability were generated on the SiC-LaB6 surface. However, owning to the loss of oxide scale in the severe ablation environment, La2Si2O7-La2O3 molten phases with their relatively low melting point could not effectively protect the ceramic. Compared with those of the SiC-LaB6 ceramic, the inside oxygen diffusion channels of the SiC-LaB6-ZrB2 ceramic were reduced owing to the improved density caused by the reactive sintering effect. Moreover, a ZrO2-La2O3-SiO2 oxide scale with good self-healing ability and excellent stability was formed during ablation, which retard further oxidation effectively. Therefore, the SiC-LaB6-ZrB2 ceramic exhibited a  better ablation resistance than the SiC-LaB6 ceramic.  Declaration of competing interest  We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled  Acknowledgements  This work was supported by the National Natural Science Foundation of China (51304249) and Natural Science Foundation of Hunan Province, China (14JJ3023).  References  [3]  [5]  [6]  [7]  [12]  [10]  [1] X. Yang, Q.Z. Huang, Y.H. 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},{
  "_id": 13,
  "PDF": "Ablation resistance of HfC(Si, O)-HfB2(Si, O) composites fabricated byone-step reactive spark plasma sintering.pdf",
  "Text": "['Journal Pre-proof  Ablation resistance of HfC(Si, O)-HfB2 (Si, O) composites fabricated by one-step reactive spark plasma sintering  Wei Hao, Na Ni, Tianyu Liu, Lei Zhou, Yao Yao, Li Ling, Juan Jiang, Yinchun Shi, Xiaofeng Zhao, Ping Xiao  PII:  DOI:  S0955-2219(20)30969-9  https://doi.org/10.1016/j.jeurceramsoc.2020.11.054  Reference:  JECS 13759  To appear in:  Journal of  the European Ceramic Society  Received Date:  14 August 2020  Revised Date:  26 November 2020  Accepted Date:  30 November 2020  Please cite this ar ticle as: Hao W, Ni N, Liu T, Zhou L, Yao Y, Ling L, Jiang J, Shi Y, Zhao X, Xiao P, Ablation resistance of HfC(Si, O)-HfB2 (Si, O) composites fabricated by one-step reactive spark plasma sintering, Journal of the European Ceramic Society (2020), doi: https://doi.org/10.1016/j.jeurceramsoc.2020.11.054  This is a PDF file of an ar ticle that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the ar ticle. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal per tain.  © 2020 Published by Elsevier.  \\x0c', 'Ablation resistance of HfC(Si, O)-HfB2(Si, O) composites fabricated by one-step   reactive spark plasma sintering   Wei Hao a, b, Na Ni a, b, *, Tianyu Liu c, Lei Zhou c, Yao Yao a, Ling Li a, Juan Jiang a, b,   Yinchun Shi a, Xiaofeng Zhao a, *, Ping Xiao a   School of Materials Science and Engineering, Shanghai Jiao Tong University,   Shanghai 200240, China   Mechanical Engineering, Shanghai Jiao Tong University, Shanghai 200240, China   c School of Materials Science and Engineering, Northwestern Polytechnical   b Key Lab of Education Ministry for Power Machinery and Engineering, School of   a Shanghai Key Laboratory of High-Temperature Materials and Precision Forming,   Journal Pre-proof  Dense HfC(Si, O)-HfB2(Si, O) composites are fabricated to enhance ablation   Zhao)   Highlights:   \\uf06c   University, Xi’an, 710072, China   * Corresponding author: Tex./Fax: +86-21-54742561   E-mail address: na.ni@sjtu.edu.cn (Na Ni) and xiaofengzhao@sjtu.edu.cn (Xiaofeng   resistance of HfC.                                                                   \\x0c', '\\uf06c   The composite is composed of HfC(Si, O), HfB2(Si, O), O-doped SiC, and free   carbon.   \\uf06c   The dense oxide layer contains Si-doped and (Si, B)-codoped HfO2, and HfB2   matrix grains.   \\uf06c   The HS15 composite sample exhibits a better ablation resistance.   Abstract:   ablation rates are the result of a slight mass gain/thickness increase, which indicates   that the oxidation process was stable and mechanical scouring was limited during   sintering using HfC and SiB6 as starting reactants. The best ablation resistance was   obtained with the composite fabricated with the addition of 15 vol.% SiB6. After   ablation under an oxyacetylene flame for 60 s, the mass and linear ablation rates of   A dense HfC(Si, O)-HfB2(Si, O) composite was fabricated by reactive spark plasma   this composite were -0.007 mg cm-2 s-1 and -0.233 μm s-1, respectively. The negative   Journal Pre-proof  Keywords: Hafnium carbide; Hafnium diboride; Solid solutions; Ablation resistance;   ablation. This enhanced ablation resistance was attributed to a unique double-layered   strength. The solid solution nature of the composite and its appropriate phase   composition were responsible for the stable oxide structure formation.   oxide formation, which possessed lower oxygen permeability and better mechanical   Solid-solution HfO2   1.   Introduction   Hafnium carbide (HfC) belongs to the class of ultra-high temperature ceramics   (UHTCs) and possesses extremely high melting point (~3900°C), high hardness and       \\x0c', 'scouring [7].   their corresponding oxides. HfO2 exhibits high melting points of ~2800°C [8, 9],   Young’s modulus, together with excellent thermo-mechanical and thermo-chemical   properties [1-3]. Therefore, it is identified as a good candidate for nose cones and   wing leading edges of hypersonic vehicles operating in extreme environments [4].   However, HfC easily oxidizes in oxygen-containing atmosphere from 400°C, forming   typically a porous HfO2 oxide layer due to the escape of gas by-products [5, 6]. This   porous oxide layer exhibits poor oxidation resistance and is vulnerable to mechanical   which can serve as skeleton for an oxide protective layer at high temperatures. HfO2   derived from the oxidation of HfC and HfB2 exhibits porous structure at middle-high   Generally, the ablation resistance of UHTCs is affected by the melting points of   Journal Pre-proof  rate in HfO2 is fast with a kp=10−7 g2 cm4 s−1 at 1500°C [9]. Glass phases such as   HfO2 framework of the oxide protective layer up to 1500 °C [5, 10-11]. In addition, kp   is only 10−11 g2 cm4 s−1 for silicates at 1500 °C [9], which can help reduce the oxygen   Si-and B-containing species have proven to be effective in improving ablation   resistance at ultra-high temperatures [4, 5, 10, 12]. A great amount of work has   diffusion through the oxide layer. Therefore, carbide-based composites incorporating   temperatures (~1800°C) [3, 4, 7]. Above 2000°C, HfO2 exhibits good mechanical   scouring resistance due to its partial densification [3], However, the oxygen diffusion   silica and borosilicate with melting point of ~1800°C can effectively fill the spaces of   focused on the MC-MB2-silicides-borides (M = Hf, Zr, Ti, and Ta) composites [4, 10,   12-14]. Young et al [4] fabricated HfC-HfB2 composites with enhanced oxidation and     \\x0c', 'ablation resistance through B2O3 filling the pores in the HfO2 layer. B2O3 volatilizes   above 1100°C and absorbs heat on the composites surface. Because the evaporation of   B2O3 requires energy, and some of the heat flux is spent for the phase transformation.   However, above 1500-1650°C, vaporization of volatile compounds leads to the   formation of pores and holes in the oxide layer which becomes vulnerable to further   volatilization of B2O3 and other gas by-products cause further oxidation and   mechanical scouring, finally decreasing the structural reliability of the B-containing   improved oxidation and ablation resistance compared with those of MoSi2/SiC   coating due to the formation of a glassy La-Si-O-B-Ti borosilicate scale providing   oxidation and mechanical scouring. These defects in the oxide layer resulting from the   composites [10]. Wang et al [12] reported a LaB6-MoSi2-TiB2 composite coating with   Journal Pre-proof  formation of a borosilicate glassy oxide layer to improve the oxidation resistance of   coatings [13, 15], and the evaporation of Si can absorb heat from the ablated sample   Incorporating free silicon into TaB2-SiC ablation resistant coating can also induce the   good protection from oxidation and mechanical peeling during ablation at ~2700°C.   The borosilicate glassy La-Si-O-B-Ti scale can effectively inhibit SiO2 and B2O3   volatilization and facilitate the formation of an integrated oxide layer. Feng et al [10]   reported the ablation resistance of B-modified SiC/HfC-SiC composite system, where   the Band Si-containing integrated oxide layer generated on the composites surface   resisted mechanical erosion and blocked oxygen diffusion into the matrix.   surface, decreasing its surface temperature [13, 14]. However, the abrupt evaporation   of gaseous free Si and other gas by-products (B2O3 (g), SiO2 (g), CO (g), and CO2 (g))   \\x0c', 'results in the depletion of Si-containing species and damage to the TaB2-SiC-Si   coating, reducing its mechanical scouring resistance at ultra-high temperatures [13].    Therefore, two principle aspects should be considered to enhance the ablation   resistance of HfC-based ceramics. Firstly, the formation of dense and continuous   oxide layers is desirable to reduce the inward oxygen diffusion and resist mechanical   two requirements obviously depends on the type and amount of the Siand B widely used in combination with MoSi2 as thermal protection coating in hypersonic   scouring. Secondly, the dense oxide layer needs to exhibit an appropriate viscosity for   the release of gas by-products so that an intact and dense structure can be maintained.   The possibility of forming an oxide scale with the favorable combination of the above   vehicles due to its high melting point (~2000°C) [16], excellent chemical stability and   containing additives, and the way they are incorporated into the matrix. SiB6 has been   Journal Pre-proof  oxidation resistance at high temperatures [17-20]. We have previously shown that   SiB6 can be employed as a unique boron and silicon source to in-situ synthesize   diborides and form HfC(Si, O) and HfB2(Si, O) solid solutions for toughening HfC   [21]. The formation of Si solid solutions in our and several other systems have been   found to improve the mechanical strength and toughness, which are due to the higher   densification degree, solid-solution strengthening, and microstructural modification   such as grain refinement [21-23]. The better mechanical performances may be   beneficial in better coping with the thermal stress and cracking during ablation [24],   finally enhancing the mechanical scouring resistance. In addition, the presence of Si   as a homogeneous solute in the HfC-HfB2 composites instead of as a secondary   \\x0c', 'separate phase may potentially modify the ablation behavior of the composites.   Silvestroni et al [25] fabricated (Zr, Mo)B2 solid solution composites, and showed that   the formation of solid solution had notable impacts on the oxidation resistance of   composites subjected to cyclic oxidation at 1650°C.   This work is therefore focused on the ablation resistance of the HfC-HfB2 solid   containing species. The microstructural evolution, sintering densification, and   ablation behavior of the composites were investigated, and the underlying ablation   solution composites, which were fabricated by in-situ reactive spark plasma sintering   mechanisms were also discussed.   2. Experimental procedure   2.1 Spark plasma sintering (SPS)   (SPS) of HfC and SiB6 powder mixtures. SiB6 served as a single additive of Si and B Journal Pre-proof  The commercially available HfC powders (purity: 99 %, main impurities include   in Fig. 1. HfC and SiB6 powders were first mixed by ball-milling with SiC balls for 8   h using isopropanol as a milling medium. Bulk samples were sintered in the presence   ZrC<0.46 %, O<0.50 %, particle size: 100 nm; Shanghai Chao Wei Nanotechnology   Co. Ltd., Shanghai, China), and SiB6 powders (purity: 98 %, main impurity includes   Si<2 %, particle size: 3-8 μm; Shanghai Alfa Aesar Co. Ltd, China) were used as the   starting materials. The XRD and SEM analyses of HfC and SiB6 powders are shown   of 0 vol.%, 10 vol.%, 15 vol.% and 20 vol.% SiB6 by SPS using a SPS furnace (FTC   HP D25, FCT Systeme GmbH, Rauenstein, Germany). Bulk samples were named as   HS0, HS10, HS15, and HS20, respectively. After drying at 60°C for 2 h, the mixture     \\x0c', 'of the above powders was put into a 31 mm graphite die which was lined with 0.3 mm   thick graphite foils to maximize the electrical and thermal conduction between the   punches and the die. The whole assembly was heated from room temperature to   1100°C at a rate of 100°C/min with an initial uniaxial pressure of 6.6 MPa, and then   from 1100°C with 6.6 MPa to 1850°C with 50 MPa at a rate of 50°C/min and 3.0   the furnace was naturally cooled to room temperature.   the applied pressure was immediately reduced to the initial pressure of 6.6 MPa, and   MPa/min under vacuum (5 Pa), respectively. After a dwell time of 20 min at 50 MPa,   Journal Pre-proof  fabricated HfC-based composites was   2.2.1 Structure characterization   The phase composition of the as‐  Fig. 1 XRD patterns (a, b) and SEM images (c, d) of the HfC and SiB6 commercial   powders: (a, c) HfC; (b, d) SiB6.   2.2 Characterization of the as-sintered HfC and HfC-HfB2 composites   analyzed by X‐  ray diffraction (XRD, Ultima IV, Rigaku, Tokyo, Japan), and Raman   spectroscopy (LabRAM HR, equipped with a focused laser spot (Nd: YAG, 532 nm)   with a diameter of about 5 μm using ×10 Objective microscopy, Horiba Jobin Yvon,       \\x0c', 'Kyoto, Japan). The microstructure and chemical composition of the HfC-based   composites were characterized by field emission scanning electron microscopy (FE SEM, MIRA3-LHM 5-20 kV, TESCAN, Brno, Czech) equipped with energy   dispersive X-ray spectroscopy (EDX, Aztec X-MaxN80 20 kV, Oxford Instruments,   Oxfordshire, UK), and scanning transmission electron microscopy ((S)TEM, TALOS   Czech). A quantitative assessment was used to estimate the phase fraction in bulk   samples by combining XRD Rietveld Refinement using the Maud software [26] and   Massachusetts, USA). TEM specimens of selected as-fabricated and ablated samples   were prepared using Focused Ion Beam (SEM-FIB, GAIA3 GMU, TESCAN, Brno,   image analysis (Image J, 1.52a, National Institutes of Health, USA) based on BSE SEM and STEM HAADF images. The true densities of the bulk samples were   F200X 200 kV, Thermo Fisher, Massachusetts, USA) equipped with energy   dispersive X-ray spectroscopy (EDX, Super-X 200 kV, Thermo Fisher,   Journal Pre-proof  The ablation resistance of HfC-based composites pellets (Φ30 × 5 mm) was   HAADF images. The grain size was evaluated and calculated for both the constituents   measured using the Archimedes’ method in distilled water. The residual porosity of   the bulk samples was evaluated by image analysis from BSE-SEM and STEM   evaluated under an oxyacetylene flame. The flame was parallel to the sample axis, and   the distance between the torch nozzle and the sample surface was 10 mm. The heat   flux during ablation was approximately 2400 kW m-2. The pressures of O2 and C2H2   of the HfC and the formed HfB2 using linear method on at least 500 grains.   2.2.2 Ablation tests   \\x0c', 'were 0.4 and 0.095 MPa, respectively. The fluxes of O2 and C2H2 were 0.244 and   0.167 L s-1, respectively. After being tested under the flame for 60 s, the samples were   naturally cooled to room temperature. During the tests, an infrared thermometer   (A600, FLIR system Inc, Oregon, USA) indicated that the highest temperature of the   central ablated surface reached approximately 2500°C. The temperature measurement   surface temperature using the same oxyacetylene flame equipment was evaluated to   instrument operating spectral range. In order to set the correct value of the spectral   emittance, the infrared thermometer measurement was compared to the actual   adjusted until the infrared thermometer gave back the same temperature as the   pyrometers, as referred in Ref. [27]. The estimated deviation of the central ablated   of the infrared thermometer is dependent on the surface emissivity ελ measured by the   temperature detected by the pyrometers in the two-color mode, and the value of ελ was   Journal Pre-proof  where Rl refers to LAR (μm s-1) and Rm to MAR (mg cm-2 s-1); l0 is the thickness of   be ~100°C (2500±100°C) in a previous work [28]. The linear ablation rate (LAR) and   the as-fabricated composites samples (μm); l1 is the sample thickness after ablation   (μm); m0 is the mass of the as-fabricated composites samples (mg); m1 is the sample   mass after ablation (mg); Δt is the ablating time (s), which comprises the entire   mass ablation rate (MAR) of the composite samples were calculated according to Eq.   (1) and Eq. (2), respectively. The reported ablation rates were the average values of   three parent specimens.    (1)    (2)   01lllRt\\uf02d\\uf03d\\uf04401mmmRSt\\uf02d\\uf03d\\uf0b4\\uf044                                                                                                                                                                                                \\x0c', 'duration of the ablation testing when the oxyacetylene flame is on the sample surface;   S is the ablated surface area of samples (cm-2).   3. Results   3.1 Structure of the HfC and HfC-based composites   Fig. 2a displays XRD patterns of the as-fabricated HfC and HfC-based   composites with different contents of SiB6 by SPS. The indexed diffraction peaks   presence of a minor HfO2 phase in the HS0 sample is attributed to a slight oxidation   of HfC during SPS at 1850°C under vacuum with an estimated oxygen partial   pressure of 1.05 Pa. Previous work has reported the oxidation of HfC and TiC over a   show that all samples primarily contain the cubic HfC phase (PDF NO. 65-8747). The   With the increase of the SiB6 content, the diffraction peaks of the hexagonal   Journal Pre-proof  HfB2 (PDF NO. 65-8678) appeared, and their peak intensities progressively increased,   carbon increased with the increase of SiB6 addition. The formed HfB2 content   increased from 34.5 vol.% in the HS10 sample to 35.2 vol.% in the HS15 sample   suggesting the in-situ formation of a HfC-HfB2 composite during the SPS process, in   agreement with our previous work [21]. The phase composition of the HfC and HfC HfB2 composites is summarized in Table 1. The contents of HfB2, O-doped SiC, and   (Table 1). Minor HfO2 peaks are still present in all of the HfC-HfB2 composites. For   wide temperature range (700-1500°C) at PO2=0.08-80 kPa [29].    the HS20 sample, the HfB2 peaks intensities are stronger than those of HfC,   suggesting that the HfB2 content (37.3 vol.%) is higher than the HfC content (28.4   vol.%). In addition, a free carbon phase (1384 cm-1 D peak, 1606 cm-1 G peak) could   \\x0c', 'be identified in all samples using Raman spectroscopy (Fig. 2b) [2, 21]. No peaks of   HfC and HfB2 were identified in the Raman spectra, which indicates that these phases   have lower Raman activity [30].   sintered by reactive SPS.   evaluated by XRD Rietveld Refinement and confirmed by image analysis.   Table 1 Phase composition of the as-fabricated HfC and HfC-HfB2 composites   Fig. 2 XRD patterns (a) and Raman spectra (b) of the HfC and HfC-based composites   Journal Pre-proof  Fig. 3 shows back-scattered electron images of the HfC and HfC-HfB2   composites sintered by reactive SPS. For the HS0 sample, small grains with a   Phase compositions   48.1   34.5   5.1   28.4   37.3   7.1   HfC   HfB2   HfO2   O-doped SiC   C   40.1   35.2   4.4   6.3   9.8   —   5.2   83.3   —   4.9   8.7   Samples   HS0   HS10   HS15   HS20   10.6   14.2   vol %   6.6   bright/gray/dark contrast can be seen (Fig. 3a), suggesting a variation in composition.   These have been confirmed by TEM characterization in our previous work [2], and   attributed to the oxygen diffusion into some of the brighter HfC grains to form       \\x0c', 'size of all phases was retained (~780 nm), and a homogeneous distribution of   HfCxOy. The dark contrast at grain boundaries is attributed to free carbon formation   during slight oxidation of HfC in SPS [2, 29]. In the HS10 sample (Fig. 3b), a   secondary phase can be observed with darker contrast amongst the brighter HfC and   gray HfB2 grains, and the grain size was evaluated to be ~750 nm (Fig. 3b and Table   2). A lower porosity of 0.8% was achieved for the HfC-HfB2 composites compared   with that of the pure HfC ceramics, around 5.7%. In the HS15 sample, a small grain   residual porosity of the HS15 sample slightly increased to 1.3% (Table 2). In the   HS20 sample, a secondary phase with large size of ~1 μm inhomogeneously   distributed amongst the HfC and HfB2 grains, and the grain size of both HfC and   HfB2 increased to ~820 nm (Fig. 3d). The residual porosity of the HS20 sample   same grain size in Fig. 3c. Free carbon was also detected at the grain boundaries. The   secondary phases with darker contrast in HfC-HfB2 composites can be observed with   Journal Pre-proof  further increased to 2.4% (Table 2). This suggests that excessive SiB6 addition can   result in reduced densification of the HfC-HfB2 composites and grain growth   acceleration.     \\x0c', 'Fig. 3 BSE-SEM images of the HfC and HfC-HfB2 composites sintered by reactive   SPS: (a) HS0; (b) HS10; (c) HS15; (d) HS20.   Table 2 Density, porosity, and grain size of the HfC and HfC-HfB2 composites   sintered by reactive SPS.   Sample   Density (g cm-3)   Porosity (%)   Grain size (nm)   HS0   10.8   HS10   11.3   9.9   9.1   reported.   680±260   567±180*   571±210*   605±217*   HS15   HS20   5.7   0.8   1.3   2.4   * The grain size of HfC and HfB2 is similar so an average value from HfC and HfB2 is   To further investigate the phase compositions and microstructure of the HfC Journal Pre-proof  contrast were identified by selected area electron diffraction (SAED) to be cubic HfC   annular dark field (HAADF) images of this sample. The grains with bright and gray   composite. Fig. 4a and Fig. 4b show typical STEM bright field (BF) and high-angle   XRD analyses. STEM-EDX mapping (Fig. 4(e-i)) suggests that the bright grains in   the Z contrast dominated STEM HAADF image contained mainly Hf, C, Si and O   (Fig. 4c) and hexagonal HfB2 (Fig. 4d), respectively, which is consistent with the   HfB2 composites, TEM and STEM-EDX analyses were carried out for the HS15   (HfC grains, Fig. 4m and Table 3), and that the gray grains contained mainly Hf, B, Si   and O (HfB2 grains, Fig. 4n and Table 3). Notably, the content of O in HfB2 grains   was higher than that of HfC grains, as shown in Fig. 4i and Table 3. These phenomena     \\x0c', 'reveal that SiB6 removed oxide impurities and reacted with HfC or HfO2 to form   HfC(Si, O) and HfB2(Si, O) solid solutions. In addition, the presence of free carbon at   the grain boundaries was also confirmed by STEM-EDX mapping analyses (Fig. 4f).   The secondary phase with dark contrast was identified by fast Fourier transform   (FFT) and inverse fast Fourier transform (IFFT) (Fig. 4j), together with TEM-EDX   point analyses to be crystalline with a cubic structure (Fig. 4(k, l)), containing Si, C   detected by XRD, presumably due to its low amount.   and O (O-doped SiC, Fig. 4o and Table 3). The O-doped SiC secondary phase was not   Journal Pre-proof  Fig. 4 STEM bright field (BF) reference image (a) and high-angle annular dark field       \\x0c', '(HAADF) image (b), SAED of HfC and HfB2 grains (c, d), STEM-EDX element   mapping (e-i), HRTEM image (j), Fast Fourier transform (FFT) pattern (k) and   inverse fast Fourier transform (IFFT) image (l) of the HRTEM image (j) of the O doped SiC phase. STEM-EDX point and area analyses (m-o) of the HS15 sample. It is   noted that the amorphous background in Fig. 4j was due to the beam damage from   FIB.   by STEM-EDX that are shown in Fig. 4b.   Table 3 Elemental compositions of different phases in the HS15 sample characterized   Journal Pre-proof  The mass and linear ablation rates of the as-fabricated HfC and HfC-HfB2   composites after ablation at 2500°C for 60 s are shown in Fig. 5. During the ablation,   HfC-based composites samples experienced oxidation, volatilization and mechanical   3.2 Microstructure of HfC and HfC-HfB2 composites upon ablation   O-doped SiC phase —   Elemental compositions   61.7   26.4 —   Hf   C   B   O   4.7   3.2   2.9   60.1 —   37.0   24.0   —   69.7   3.1   HfC grain   HfB2 grain   Different phases   at %   Si   7.2   scouring simultaneously. The first process causes mass gain/thickness increase while   the latter two result in mass loss/thickness reduction [31]. The mass and linear         \\x0c', 'ablation rates reflect the combined effect of oxidation, volatilization and mechanical   scouring. For the HS0 sample, the mass and linear ablation rates of HfC ceramic were   0.031 mg cm-2 s-1 and 0.733 μm s-1, respectively. With increasing contents of SiB6   addition, the mass and linear ablation rates gradually reduced. In the HS15 sample,   the mass and linear ablation rates were -0.007 mg cm-2 s-1 and -0.233 μm s-1,   place during the ablation process.   respectively, which demonstrated that both mass gain and oxide thickening were   promoted thanks to the formation of a dense oxide layer. This inhibited the weight   loss from mechanical erosion. In the HS20 sample, the composite sample enabled   large mass gain and increasing thickening, which indicated that severe oxidation took   Journal Pre-proof  Surface XRD patterns and Raman spectra of the HfC and HfC-HfB2 composites   after ablation are shown in Fig. 6. In all of the ablated samples, the oxide layer on the   surface was composed of monoclinic HfO2 as result of the reactions of Eq. (3-5) [32].   SEM surface images of the as-fabricated HfC and HfC-HfB2 composites after ablation   are displayed in Fig. 7. For the HS0 sample, some buckling was generated with   Fig. 5 Ablation indicators of the HfC and HfC-HfB2 composites at 2500°C for 60 s:   (a) mass ablation rate; (b) linear ablation rate.       \\x0c', 'formation of microholes on the surface of the oxide layer, and microcracks can also be   observed (Fig. 7a and Fig. 7e). In the HS10 sample, the number and size of   microholes increased (Fig. 7b and Fig. 7f), but the microcracks disappeared. Notably,   in the HS15 sample, a dense and homogeneous oxide layer formed with few   microholes (Fig. 7c and Fig. 7g). In the HS20 sample, the amount and size of   microholes on the sample surface increased again (Fig. 7d and Fig. 7h).   Fig. 6 Surface XRD patterns (a) and Raman spectra (b) of the HfC and HfC-HfB2   Journal Pre-proof  composites after ablation.    (3)    (4)    (5)   Fig. 7 SEM surface images of the HfC and HfC-HfB2 composites after ablation: (a, e)   HS0; (b, f) HS10; (c, g) HS15; (d, h) HS20.   ()2 ()2()()2HfC3O2HfO2COsgsg\\uf02b\\uf0ae\\uf02b()2 ()2()2()HfC2OHfOCOsgsg\\uf02b\\uf0ae\\uf02b2()2 ()2()23()2HfB5O2HfO2BOsgsl\\uf02b\\uf0ae\\uf02b                                                                                                                                                              \\x0c', 'SEM cross-sectional images of the HfC and HfC-HfB2 composites after ablation   are shown in Fig. 8. For the HS0 sample, the HfO2 oxide scale in the HfC sample   exhibited a double-layered structure, containing a dense upper layer and a porous   bottom layer (Fig. 8a and Fig. 8e). An oxidation diffusion layer (HfCxOy) with   thickness of ~70 μm was observed underneath the double-layered oxide. The   thickness of the whole oxide layer was 131±15 μm (Fig. 9). These observations are   contained in the oxide layer, and the contents of Si and B gradually decreased from   the oxide layer to the HfC-based composite matrix (Fig. 8d and Fig. 8l). It is noted   that the oxygen diffusion depth into the matrix from the interface in the HS15   dense, and homogeneous oxide layer formed and exhibited a double-layered structure   consistent with our previous work [7] and literature data [3]. In the HS10 sample,   similar horizontal microcracks at the interface between the oxide layer and the   ceramic matrix were observed. The oxide thickness decreased to 75±10 μm in the   HS10 sample (Fig. 9). Significantly, In the HS15 sample, a completely crack-free,   with thickness of 64±6 μm (Fig. 9), including an upper layer and a dense bottom layer   Journal Pre-proof  this sample (Fig. 8j). Also in the HS20 sample, the double-layered structure oxide was   and the upper/bottom oxide layer interface, respectively. Fig. 8l includes the EDX   analyses of the HS20 sample from the areas marked in Fig. 8d. Si and B were also   (Fig. 8c and Fig. 8g), which were further confirmed with high magnification SEM   images (Fig. 8j and Fig. 8k). A few micropores were observed in the upper layer of   confirmed with a thickness of 65±8 μm (Fig. 8d, Fig. 8h and Fig. 9). However, large   penetrating vertical and horizontal cracks were also observed in the upper oxide layer   \\x0c', 'composites matrix (~20 μm) was lower than that in the HS20 composites matrix (~44   μm away), confirming that the oxide scale formed on the former sample had a better   performance as oxygen barrier.   Journal Pre-proof  HfB2 composites after ablation: (a, e) HS0; (b, f) HS10; (c, g, i-k) HS15; (d, h) HS20;   Fig. 8 SEM (a-d) and BSE-SEM (e-h) cross-sectional images of the HfC and HfC (l) elemental compositions of different areas in the oxide layer of HS20 sample after   ablation at 2500°C that are shown in Fig. 8d.         \\x0c', 'Fig. 9 Thickness of oxide layer as a function of the HfB2 content for the HfC and   HfC-HfB2 composites after ablation.   To investigate the composition of different areas in the oxide layers in the HS10   composite after ablation, SEM-EDX area analyses and elemental mapping were   performed, as shown in Fig. 10. The upper layer contained Hf, O, C, and B (“Area Ⅰ”   in Fig. 10a and Fig. 10g). The bottom oxide layer exhibited a denser microstructure   than the upper layer, and contained Hf, O, C, and B too (“Area Ⅱ” in Fig. 10a and   Fig. 10g). The content of B in the dense bottom layer was higher than that of the   upper layer (Fig. 10g). The unoxidized HfC-HfB2 matrix primarily contained Hf, C,   Raman shifts corresponding to those of monoclinic HfO2 phase [32], which confirmed   B, and Si with a small amount of O (“Area Ⅲ” in Fig. 10a and Fig. 10g). Raman   spectra from the upper (Area Ⅰ) and bottom (Area Ⅱ) oxide layers (Fig. 11) showed   Journal Pre-proof  that both of oxide layers primarily contained HfO2. Raman data also revealed that the   HfC-HfB2 composite matrix (Area Ⅲ) contained a few HfO2 phase.       \\x0c', 'Fig. 10 SEM cross-sectional image (a) and SEM-EDX element mapping analyses (b f) of the HS15 sample after ablation, (g) elemental compositions of different areas   indicated in Fig. 10a.   ablation at 2500°C that are shown in Fig. 10a.   STEM and STEM-EDX analyses. Fig. 12a and Fig. 12b show STEM BF and STEM   HAADF images of the upper oxide layer, respectively. Comparatively, these results   revealed that the upper oxide layer had a few micropores (Fig. 12a), which is   Fig. 11 Raman spectra of different areas of the oxide layer in the HS15 sample after   The upper oxide layer in the HS15 composites was further characterized by   Journal Pre-proof  have been observed in previous works [33, 34]. The areas with the dark contrast were   consistent with the SEM observations (Fig. 8j). HRTEM, SAED, and STEM-EDX   area analyses (squared “Area (Ⅰ-Ⅲ)” in Fig. 12b, Fig. 12(c-e) and Fig. 12(k, l))   together suggested that a large amount of Si-doped and (Si, B)-codoped monoclinic   HfO2 grains formed in the upper oxide layer. Both of the above two solid solutions   found by STEM-EDX mapping analysis to contain only C (Fig. 12e and Fig. 12(f-j)),   which confirmed the depletion of O-doped SiC phase in the upper oxide layer. Some   of the smaller oxide grains in the upper oxide layer contained Hf, O and Si, which     \\x0c', 'may be a Si-doped HfO2 phase (squared “Area Ⅳ” in Fig. 12e, Fig. 12m), suggesting   the Si-doped and (Si, B)-codoped HfO2 composite scale formation through oxidation   of HfC(Si, O) and HfB2(Si, O) solid solutions, and O-doped SiC.   To summarize, the upper oxide layer of the HS15 sample mainly comprised of   Si-doped and (Si, B)-codoped monoclinic HfO2 grains as well as residual carbon and   a few micropores.   Journal Pre-proof  Fig. 13 shows the TEM characterization of the dense bottom layer of the HS15   Fig. 12 TEM analysis of the upper oxide layer in the HS15 sample after ablation: (a)   mapping analyses; (k-m) STEM-EDX spectra (squared areas in (e)).   STEM bright-field (BF) reference image; (b) STEM high angle annular dark-field   (HAADF) image; (c, d) HRTEM image and SAED pattern of HfO2 grain (squared   “Area Ⅰ” in (b); (e, f-j) STEM HAADF image and corresponding STEM-EDX   sample. It can be observed that the denser oxide layer had fewer micropores compared   with the upper oxide layer (Fig. 13a and Fig. 13b). The bottom oxide layer was   mainly composed of gray/bright monoclinic HfO2 grains, as confirmed by SAED       \\x0c', 'analysis (squared area in Fig. 13a, Fig. 13(b, c)). The gray monoclinic HfO2 grains   contained Hf, O, B and Si, as revealed by the STEM-EDX elemental mapping (Fig.   13(e-h)) and STEM-EDX spectra analysis (squared “Area Ⅰ” in Fig. 13b, Fig. 13i),   while the bright monoclinic HfO2 grains only contained Hf, O, and Si (squared “Area   Ⅱ” in Fig. 13b, Fig. 13j). These results indicated similar results as before, i.e. the   bright/gray monoclinic HfO2 grains were Si-doped and (Si, B)-codoped monoclinic   TEM characterization suggested that the dense bottom oxide layer of the HS15   sample was mainly composed of Si-doped and (Si, B)-codoped monoclinic HfO2   HfO2 grains, respectively. Significantly, some of HfB2 matrix with hexagonal   structure were embedded in the Si-doped and (Si, B)-codoped HfO2 grains, as   revealed by SAED (Fig. 13a and Fig. 13d). The grain boundary phases only contained   Hf and O, confirming the presence of HfO2 (squared “Area Ⅲ” in Fig. 13b, Fig. 13k).   Journal Pre-proof  Furthermore, free carbon was not found in this dense bottom oxide layer. Overall, the   grains, and some unoxidized HfB2 grains.   Fig. 13 TEM analysis of the bottom oxide layer in the HS15 sample after ablation: (a)     \\x0c', 'STEM BF reference image; (b) STEM HAADF image; (c, d) SAED patterns of HfO2   and HfB2 grains (squared areas in (a)); (e-h) STEM-EDX mapping analyses; (i-k)   STEM-EDX spectra (squared areas in (b)).   4. Discussion    4.1 Effect of SiB6 addition on the microstructure evolution of the HfC-HfB2   composites   reactions as shown in Eq. (6), Eq. (7) and Eq. (8) are proposed to occur during the   reaction was calculated according to Eq. (9) [21].   SPS process. The change in Gibbs free energies (ΔGT/kJ mol−1, T=1850°C) for each   Based on the comprehensive XRD, Raman and TEM data, the following   =-175.155 kJ mol-1)                                      (6)   Journal Pre-proof  =-672.006 kJ mol-1)                                      (7)    is the Gibbs formation free energy for each compound, kJ mol-1;   ,   =-48.522 kJ mol-1)       (  (  (  where    (8)    (9)   is the coefficient of reactants or products;    is the change in Gibbs free energy of   the reaction at 1850°C under vacuum atmosphere (5 Pa), kJ mol-1; R is the ideal gas   constant, J mol-1 K-1; T is temperature, K;    is the gas product pressure, Pa;    is   ()6()2()()()3HfC + SiB3HfBSi3Csssls\\uf0ae\\uf02b\\uf02bTG\\uf0442()6()()2()()(g)3HfOSiB6C  3HfBSi 6CO  ssssl\\uf02b\\uf02b\\uf0ae\\uf02b\\uf02b\\uf0adTG\\uf044()(s)()SiCSiCls\\uf02b\\uf0aeTG\\uf044()lnfrffTifTirrppGGiRTpp\\uf06e\\uf071\\uf071\\uf06e\\uf071\\uf06e\\uf0e9\\uf0f9\\uf0e6\\uf0f6\\uf0ea\\uf0fa\\uf0e7\\uf0f7\\uf0ea\\uf0fa\\uf0e8\\uf0f8\\uf044\\uf03d\\uf044\\uf02b\\uf0ea\\uf0fa\\uf0e6\\uf0f6\\uf0ea\\uf0fa\\uf0e7\\uf0f7\\uf0ea\\uf0fa\\uf0e8\\uf0f8\\uf0eb\\uf0fb\\uf0d5\\uf0e5\\uf0d5()fTGi\\uf071\\uf044i\\uf06eTG\\uf044fprp                                                                                                                                                                              \\x0c', 'oxides in the HfC matrix and grain boundaries retarded the grain boundary diffusion   and mass transfer during SPS process, which resulted in a lower degree of   the gas reactant pressure, Pa;    is the coefficient of reactants;    is the coefficient   of products. The change in Gibbs free energy of all reactions is negative, indicating   that the proposed reactions are thermodynamically favorable.    The addition of SiB6 had a favorable effect on the densification and   microstructural evolution of the HfC-HfB2 composites. For the HS0 sample, it was   composed of cubic 83.3 vol.% HfC, 4.4 vol.% free carbon, and a small amount of 6.6   phase was observed, in agreement with our previous work [21]. The reaction amongst   densification, and a residual porosity of 5.7%. In the HS10 sample, the formation of   HfC(Si, O) and HfB2(Si, O) solid solutions with the presence of the O-doped SiC   vol.% HfO2 (Table 1) due to the oxidation of a small amount of HfC. The presence of   Journal Pre-proof  containing Hf, O, Si, C, and B during SPS at 1850°C [35]. The Si impurity from SiB6   low melting point [35]. HfC(Si, O) and HfB2(Si, O) solid solutions may have grown   by epitaxial precipitation from the transient Hf-O-Si-C-B liquid [36-38]. At the same   molten phase [39]. A synergic effect of diffusion mass transfer via Si solid-solutions   [2, 40] and reactive sintering contributed to the enhanced densification of the HfC HfB2 composites, with porosity from 5.7% to 0.8%. The addition of SiB6 was also   found to be effective in removing the oxides at the grain boundaries through reaction   SiB6, carbon, HfC, and HfO2 resulted in the formation of a transient liquid phase   raw powders may also have promoted the formation Hf-O-Si-C-B liquid due to its   time, the new O-doped SiC phase was also formed by a precipitation from the Si-O-C   r\\uf06ef\\uf06e  \\x0c', 'of Eq. (7), which facilitated the grain boundary diffusion and densification [21]. In   addition, grain refinement was observed in the composite samples (Table 2), which   was considered the result of a reactive process [41], leading to the formation of   HfC(Si, O) and HfB2(Si, O) solid solutions as well as of Zener pinning of the   secondary phases and free carbon at the grain boundaries [2, 40].    diffusion and inhibited the densification (Table 2) [2, 21]. It is noted that the size of   small micropores may have increased due to the preparation of TEM sample by FIB.   SPS. However, some micropores were generated at the O-doped SiC/HfC and O doped SiC/HfB2 interfaces, as revealed by the STEM DF images and STEM-EDX   mapping analyses (Fig. 14). The formation of micropores was attributed to the   presence of O-doped SiC and free carbon, which retarded the grain boundary   vol.% O-doped SiC phases were homogeneously produced through the in-situ reactive   Furthermore, for the HS15 sample, a large amount of 35.2 vol.% HfB2 and 8.7   Journal Pre-proof  Fig. 14 STEM DF image (a) and STEM-EDX mapping analyses (b-f) of the HS15   sample.       \\x0c', 'In the HS20 sample, HfB2 was the major phase, 37.3 vol.%, and the content of   HfC was only 28.4 vol.% (Table 1). Simultaneously, some grains exhibited abnormal   grain growth during in-situ reactive SPS, resulting in an inhomogeneous distribution   of grain size. The secondary O-doped SiC phase and free carbon were also seen to be   more localized and coarser with the highest contents of 10.6 vol.% and 14.2 vol.%,   below.   4.2 Ablation mechanisms of the HfC-HfB2 composites   densification, resulting in the higher porosity (2.4%) of this sample (Table 2).   respectively (Table 1). The inhomogeneous distribution appeared to inhibit the   chemical analyses, the possible reactions involved during the ablation of the HfC(Si,   O)-HfB2(Si, O) composites containing O-doped SiC and free carbon are proposed as   Based on the literature [6-7, 10, 12-13] and the above microstructure and   Journal Pre-proof   (10)    (11)    (12)    (13)    (14)    (15)    (16)    (17)   Comparatively, the HS0 sample exhibited the highest ablation rate due to severe   oxidation and mechanical scouring (Fig. 5). Buckling, microcracks, and micropores   ()2()2()2()2SiOC3O2SiO2COsglg\\uf02b\\uf0ae\\uf02b()2()2()()SiOCOSiO2COsglg\\uf02b\\uf0ae\\uf02b2()2()23()2223()HfO+ SiO+ BOHfOSiOBOslllxyzxyz\\uf0ae23()23()BOBOlg\\uf0ae2()2()SiOSiOlg\\uf0ae()2()2()C + OCOsgg\\uf0ae()2()()2C + O2COsgg\\uf0ae()2()2()()HfC  + OHfOC sgss\\uf0ae\\uf02b                                                                                                                                                                                                                                                                                                                                                                                                                                                                                           \\x0c', 'no silica, B2O3 or borosilicate glass was found in the sample after ablation, these   phases were still expected to form during ablation according to Eq. (5) and Eq. (10 formation on the surface of the ablated samples are attributed to mechanical scouring   and volume expansion induced by rapid oxidation (Fig. 7(a, e)) (Eq. (3-4)) [4, 7]. The   typical double-layered oxide layer formed on the HfC sample [3]. Penetrating   horizontal cracks were observed at the interface between upper oxide layer and   diffusion oxidation layer (Fig. 8(a, e)) due to the thermal mismatch stress (coefficients   of thermal expansion: HfC: 6.7 × 10-6 K-1; HfO2: 5.8 × 10-6 K-1) [9, 42]. These results   12). Their formation could have led to a softening of the oxide layer, which reduced   The HS20 sample therefore exhibited the poorest ablation resistance. Although   Journal Pre-proof  and B2O3(g) from the reactions described by Eq. (3-5), Eq. (10-11) and Eq. (13-16)   during ablation, which can result in some microholes formation in the oxide layer   and therefore would have much lower CTE than the upper oxide layer, leading to a   large thermal mismatch gap between the upper and bottom oxide layer [43], which   obtained with a larger amount of O-doped SiC, 10.6 vol.% and carbon, 14.2 vol.%.   They together resulted in the formation of more gas by-products (CO, CO2, SiO2(g)   its resistance to mechanical scouring. Higher HfB2 content of 37.3 vol.% was   have been also shown in our previous work [7].   (Fig. 7(d, h)) [12, 13]. The bottom oxide layer contained more B-containing oxides   can induce tensile stress in the upper layer and the formation of vertical cracks (Fig.   8(d, h)). The vertical cracks deflected at the interface between the upper and bottom   oxide layer and resulted in horizontal cracking, finally leading to the oxide layer     \\x0c', 'delamination (Fig. 8d) [44, 45].   Significantly, the HS15 sample exhibited good ablation resistance at 2500°C.   Based on the experimental results, the overall ablation mechanism of the HS15 solid   solution composite is illustrated in Fig. 15. A dense double-layered oxide structure   completely formed in this system, to which a couple of factors may have contributed.   Firstly, the phase diagram of the Hf-Si-O system demonstrated that the liquid-solid   phase transformation temperatures ranged from ~1800 to ~2000°C [35]. It appears   that the obtained composition of 40.1 vol.% HfC, 35.2 vol.% HfB2, 8.7 vol. % O doped SiC in this sample had facilitated the formation of an oxide layer with optimal   HfC(Si, O) and HfB2(Si, O) seems to have promoted the formation Si-doped and (Si,   products and sealing the oxide layer. Secondly, the solid solution nature of the starting   Journal Pre-proof  B)-codoped HfO2 upon ablation instead of simple HfO2 as reported in many previous   studies [3, 10, 12]. The reaction processes can be described by Eq. (3-5), Eq. (10-12),   during the ablation process, limiting the fierce destructive volatilization of Si   containing species [13] to help maintaining the oxide integrity. Furthermore, the   oxidation of Si as a solute was also expected to be beneficial in forming more   homogeneously distributed Si-containing phases, thereby reducing turbulence or   viscosity at the ablation temperature, which could be beneficial for releasing gas by and Eq. (17). The formation of Si-doped HfO2 scale was shown to be   thermodynamically favorable [33, 34] and suggested that Si has been better retained   bursting phenomena on the glass layer surface due to volatilization, as suggested by a   previous work on the oxidation of (Zr, Mo)B2 solid solution composites at 1650°C   \\x0c', '[25].   Fig. 15 Schematic illustration of the ablation mechanisms of the HS15 composite   sample.   Consequently, the capability of forming a dense protective double-layered oxide   Journal Pre-proof  all, the good oxidation diffusion barrier performance of the double-layered oxide was   could effectively reduce oxygen diffusion [12]. Especially, these Hf-Si-O and Hf-Si can provide advantages for the enhanced ablation resistance in two aspects. First of   evident, and demonstrated by the formation of free carbon at the oxide grain   boundaries (Fig. 12e and Fig. 12(f-j), Eq. (17)) [7, 29], the presence of both HfB2   grains in the bottom oxide layer, and the intermediate HfCxOy layer formation (Fig.   15) [6]. This result demonstrated that the Siand B-containing upper oxide layer   O-B oxide layers exhibited good oxidation resistance due to their dense structure and   possibly low oxygen permeability [7, 46, 47]. Secondly, a good mechanical strength   to resist mechanical scouring was expected due to the dense structure of the oxide     \\x0c', 'layer.   5. Conclusions    HfC-HfB2 composites containing HfC(Si, O) and HfB2(Si, O) solid solutions, O doped SiC, and free carbon were successfully fabricated starting from HfC and 0-20   vol.% SiB6 by one-step in-situ reactive SPS at 1850°C. The following conclusions can   be drawn:   compared to the pure HfC having a porosity of 5.7%.   low porosity of 1.3% and the best ablation resistance measured at 2500°C in an   oxyacetylene flame.   with 10 vol.% SiB6, which exhibited a significant densification improvement as   (i) Dense HfC-HfB2 composites with the lowest porosity of 0.8% were obtained   (ii) The addition of 15 vol.% of SiB6 resulted in the HS15 composite displaying a   Journal Pre-proof  sample, including HfC(Si, O) and HfB2(Si, O) solid solutions, O-doped SiC, and free   carbon, and its appropriate phase composition were believed to be responsible for the   (iii) The enhanced ablation resistance was attributed to the formation of a unique   double-layered oxide structure, which possibly possessed lower oxygen permeability   and better mechanical scouring resistance. The solid solution nature of the HS15   formation of a stable oxide structure.   Declaration of interests   ☒ The authors declare that they have no known competing financial interests or personal               \\x0c', 'relationships that could have appeared to influence the work reported in this paper.   Acknowledgements   This work was supported by the National Natural Science Foundation of China   (No. U19A2099, 52072238 and 51902197) and the Shanghai Pujiang Program (No.   Research Center, School of Materials Science and Engineering, Northwestern   Polytechnical University.   18PJ1406500). Ablation tests were performed in the Carbon/Carbon Composites   Journal Pre-proof          \\x0c', 'References   [1] L. Charpentier, M.B. Pichelin, J.L. Sansa, D. Scitib, L. Silvestroni, Effect of high   temperature oxidation on the radiative properties of HfC-based ceramics, Corros.   Sci. 126 (2017) 255-264.   [2] W. Hao, N. Ni, Y. Guo, C.W. Li, X.H. Fan, W.W. Xiao, X.F. Zhao, P. Xiao,   [3] C. Zhang, B. Boesl, A. 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},{
  "_id": 14,
  "PDF": "Ablation tests on HfC- and TaC-based ceramics for aeropropulsive applications.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 35 (2015) 1401-1411  Ablation tests on HfCand TaC-based ceramics for aeropropulsive applications  L. Pienti a , D. Sciti a,∗  , L. Silvestroni a , A. Cecere b , R. Savino b  a CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy b Department of Industrial Engineering (DII) - Aerospace section, University of Naples “Federico II”, Naples, Italy  Received 18 September 2014; received in revised form 12 November 2014; accepted 17 November 2014  Available online 24 December 2014  Abstract     HfCand TaC-based composite materials containing 15 vol% MoSi2 as sintering aid were selected for their excellent combination of ceramic and metallic properties to produce simple shaped prototypes for ablation tests in a mixture of oxygen/butane/propane at temperatures between 1300 C and 1900 C. The ablation resistance of  the HfCand TaC-based models was compared  to  that of graphite models. SEM analysis of  the sample surfaces and sections was conducted to analyze the material degradation. The composites have a superior ablation resistance compared to graphite, showing only slight surface oxidation and weight variation for temperatures above 1700 C. © 2014 Elsevier Ltd. All rights reserved.        Keywords: Hafnium carbide; Tantalum carbide; Hot corrosion; Oxidation; Microstructure  1.   Introduction     The thermal, chemical, and mechanical environments typical of aero-propulsion applications, such as  those characteristic of combustion chambers or of high performance  rocket nozzles, introduce many problems from  the point of view of materials. These environments are  typically characterized by highly corrosive atmospheres that may also contain metal additives, with flame temperatures overpassing 3000 The materials used  for  these  applications  include  refractory metals,  refractory-metal carbides, graphite, ceramics and fiber-reinforced composites.2,4-9 Certain classes of materials demonstrated  superior performances under  specific operating conditions, but  the choice depends on  the specific application. For  instance,  fully densified  refractory-metal nozzles generally are more  resistant  to erosion and  thermal-stress cracking than  the other materials. Graphite performs well with  the  least oxidizing propellants, otherwise  is  severely eroded. Some of the refractory-metal carbide nozzles show outstanding erosion  C.1-6  ∗  Corresponding author: Tel.: +39 0546 699723; fax: +39 0546 46381.  E-mail address: laura.pienti@istec.cnr.it (L. Pienti).  http://dx.doi.org/10.1016/j.jeurceramsoc.2014.11.018  0955-2219/© 2014 Elsevier Ltd. All rights reserved.  resistance, comparable to that of the best refractory-metal materials, but generally  suffer  from  fractures  induced by  thermal stresses.2 The interaction of environmental conditions and the dimensional  constraints  in  the  nozzle  throat makes  the  selection of  suitable  rocket nozzle materials  extremely difficult. This work  is  focused on  the development and  testing of ceramics based on carbides of early  transition metals, e.g. hafnium and tantalum carbides. These compounds possess an excellent combination of properties  including extremely high melting point, high electrical and  thermal conductivity, good  thermal shock resistance and  superior ablation  resistance compared  to C/C composites.4,10 For  the above  reasons  they have been already identified as excellent candidates  for aerospace applications, including the possibility to develop thermal protection systems for hypersonic atmospheric  re-entry conditions,  together with other ultra-refractory ceramic composites.2-4,6,8,10 As a consequence of  their extreme refractoriness, fully dense materials can be achieved only if very high temperatures and mechanical pressure are applied.  In  recent years however,  there has been a strong effort  to process HfC and TaC with different sintering  techniques and use of efficient  sintering aids.11-20 Some of  the authors of  the present work  suggested  the addition of        \\x0c', '1402   L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411  transition metal  silicides.12-15 MoSi2 was  found  to be most effective  sintering agent, allowing densification even without external mechanical pressure.15 The long-term purpose of this program is to develop ablationresistant materials  that could be reliably used with high-flame temperature rocket propellants for aerospace propulsion applications. Usually, materials for typical rocket nozzles are designed into suitable configurations only after preliminary characterization and test firings in relevant conditions of flame temperature, combustion products, gas velocity, etc. This work  is dedicated  to preliminary ablation  tests of bulk hafnium and tantalum carbides-based ceramics at relatively high temperature  in  significant combustion environments. For  the sake of comparison, pyrolytic graphite prototypes were also machined and  tested with  the same experimental setup. Additionally,  fabrication procedures,  room  and high  temperature properties and oxidation resistance of the ceramics are presented and discussed.  2. Experimental details  2.1. Materials preparation  The composites were produced starting from commercial raw materials:  •  • HfC (Cerac Inc., Milwaukee, WI, USA), particle size range 0.2-1.5  \\u242em, O: 0.77 wt%;  TaC (Treibacher Industrie AG, Althofen, Austria), mean particle size 1.1  \\u242em, O: 0.01 wt%; \\u242em, Aldrich, Steinbeim, Germany), particle size  MoSi2 (<2  1 wt%. range 0.3-5  \\u242em, O:   •  The following compositions, Table 1, were prepared (vol%) on the basis of previous studies12,18 :  • HfC + 15 vol% MoSi2 (labeled as HCM);  TaC + 15 vol% MoSi2 (labeled as TCM).  •  The powders were dispersed in absolute ethanol by simultaneous application of magnetic stirring and ultrasonication. The powder mixtures were  further ball milled  in absolute ethanol for 24 h with SiC milling media. After milling, the slurries were dried  in a rotary evaporator, sieved and shaped as 3 cm diameter discs. The samples were hot pressed at  temperature  in  the range 1800-1900 C in low vacuum (100 Pa) using an inductionheated graphite die with a uniaxial pressure of 30 MPa, following     the same schedule of previous studies.12 After sintering, the bulk densities were measured by the Archimedes’ method.  2.2. Microstructural and thermo-mechanical characterization, oxidation tests  ×  ×  ×  ×  The microstructures were  analyzed  using  scanning  electron microscope  (FE-SEM, Carl Zeiss Sigma NTS GmbH, Oberkochen, Germany)  and  energy-dispersive  spectroscopy (EDS, X-Act, INCA Energy 300, Oxford Instruments, Abingdon, UK) on polished and fractured surfaces. Mean matrix grain size, amount of porosity and amount of secondary phases were determined on micrographs of polished sections using  image analysis  (Image-Pro Analyzer 7.0, Media Cybernetics, Silver Spring, MD, Rockville, USA).  In particular,  the mean grain size of  the matrix grains was calculated by  the circle method on  the polished section, at  least 100 grains per specimen were measured.21 As  for  the mechanical properties, some values were  taken from previous  studies.12 Just  for  the  sake of clarity, Vickers micro-hardness, HV1.0, was measured with a  load of 9.81 N using a  standard micro-hardness  tester  (Zwick 3212, Zwick, Ulm, Germany), Young’s modulus, E, was measured by  the resonance frequency method on 28 mm   8 mm   0.8 mm specimens using an HP gain-phase analyzer. Fracture toughness, KIc , was evaluated using chevron-notched beams (CNB) in flexure, on bars 25 mm   2 mm   2.5 mm (length by width by thickness, respectively) using a universal machine  (Instron mod. 6025, Instron, High Wycombe, UK). On  the same machine and with the same fixture,  the 4-pt flexural strength,  σ , was measured from  room  temperature up  to 1200 C  in argon flux on  test bars 25 mm   2 mm   2.5 mm  (length by  thickness by width, respectively), using a crosshead speed of 0.5 mm/min. The  following  thermal properties were measured  in order to  complete  the  set  of measurements:  thermal  expansion coefficient, CTE,  up  to  1300 C  under flowing  argon with a  5 C/min  heating  rate,  using  a  dilatometer Netzsch mod. DIL E  402  (Netzsch, Geraetebau, Germany),  on  test  bars 25 mm   2.5 mm   2 mm (length by thickness by width, respectively);  thermal  conductivity, KTH ,  at different  temperatures from RT to 1500 or 1900 C was calculated using the following expression: KTH = DTHCPρ, where DTH is the thermal diffusivity, CP is  the specific heat and  is  the sample density. DTH was measured using a ‘laser flash’ equipment (mod. 427 Netzsch Geratebau, Germany) in argon on 10 mm   10 mm   3 mm specimens, CP was measured using a dedicated differential scanning calorimetry (2920 MDSC, TA Instruments, USA). Preliminary oxidation  tests on regular samples with dimensions 10-13 mm   2.5 mm   2 mm were carried out  for 5 min  ×  ×  ×  ×  ×  ×  ×     ρ        ×     Table 1  Labels, sintering conditions (TMAX , dwell time, pressure applied), experimental and relative densities and mean grain size.  Label   HCM   TCM   Compositions   HfC + 15MoSi2 TaC + 15MoSi2  Sintering conditions (     1900, 10, 30   1800, 3, 30   C, min, MPa)   Exp. density (g/cm3 )   Rel. density (%)   m.g.s. (\\u242em)  12.03   12.70   99.8   99.5   2.4   1.2   ± ±   1.6   0.6    \\x0c', 'L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411   1403  Fig. 1.   (a) View of the test sample, (b) drawing (mm).     in a bottom  loading furnace box when  the maximum  temperature of 1600 C was reached. Before the test, all the specimens were cleaned in solvents and their dimensions (length, thickness and width) measured. The extent of degradation after oxidation was evaluated by calculating  the percentage of mass variation,  mi )/mi , where mi and mf represent the initial and final weights,  respectively. The microstructural modifications induced by oxidation were evaluated by X-ray diffraction (Bruker D8 Advance, Bruker, Karlsruhe, Germany) and SEMEDS analysis on the surfaces.   (mf −  \\x01m = 100   ×  2.3. Ablation tests in combustion flame environment  An experimental setup was developed at the Laboratory of the University of Naples  to  test composite materials under  typical conditions of the propulsion system. Combustion flame of oxygen and hydrocarbon gases (butane-propane) was used, with an oxygen/fuel ratio of 3.5 (see Table 2). A simple model with cylindrical geometry was machined from sintered pellets (Fig. 1a). Fig. 1b shows the dimensions of the model. For the sake of comparison, commercial graphite (Tokai Carbon G 330) was used to produce models with  the same shape and dimensions of  the ceramic ones. Each sample was exposed for 4 min at different flames intensity pointing  the whole  top surface  (diameter = 1 cm, Fig. 1). Fig. 2 shows  the experimental setup utilized for each  test. The test sample (1), was mounted on an alumina support (2), able to thermally insulate the sample and to hold out at high temperature environment. A secondary ceramic support (3) was utilized  to fix the sample at given distance from the jet flame (4). Diagnostic system  included a FLIR ThermaCAM P40  thermal camera  Fig. 2. Experimental setup: (1) test sample, (2) alumina support, (3) ceramic sup port, (4) burner, (5) FLIR ThermaCAM P40 thermal camera, (6) MikronImpac  ISQ5 pyrometer, and (7) CCD camera.        (5) and a dual color MikronImpac  ISQ5 pyrometer  (6) which allows high accuracy  temperature measurements  in  the  range from 600 C to 3000 C. In addition, a CCD camera (7) was utilized  to acquire movie of each  test. Frontal face of  test sample was directly exposed to the burner flame. The distance was fixed in order  to have  the maximum  temperature of  the flame. Both thermal and pyrometric systems were pointing toward the same area of the lateral surface. Thermal images and pyrometer data were acquired and stored on a working station during each  test. Post processing of combined data allowed retrieving the thermal map of the sample and the emissivity of  the materials at  the wavelength of pyrometer (1  \\u242em) and of  the  thermal camera  (8-10  \\u242em).  In addition,  in order to measure the mass loss due to the ablation phenomena, the  samples were weighed before and after  the  test using a high precision Mettler Toledo AT261 Delta Range microbal±10  ance with an accuracy of  \\u242eg. The extent of degradation was evaluated by calculating  the percentage of mass variation,  mi )/mi , where mi and mf represent the initial and final weights and by direct measure of  the models dimensions before and after the ablative tests, respectively.   (mf −  \\x01m = 100   ×  Table 2  Conditions and results of the ablative tests, exposure time was 4 min.  Test  Sample  Flow rate (g/s)   Emissivity   Max temperature (     C)   \\x01m (%)   \\x01h center (mm)   \\x01h side (mm)   \\x01Ø (mm)  Oxygen   0.035   0.035   0.035   0.070   0.070   0.070   0.070   Fuel   0.01   0.01   0.01   0.02   0.02   0.02   0.02   Graphite   HCM   TCM   Graphite   HCM   TCM   TCM   at 1   \\u242em   at 8-10   \\u242em   1   -   -   1   -   -   -   0.85   0.65   0.7-0.8   0.85   0.65   0.85   0.85   1327   1527   1627   1627   1927   1685   1727   1  2  −13.88  −0.11  +0.17  −26.21  +0.15   +0.53  −0.23   9.46   −0.13  -  −0.08  −0.02  -  −0.07  -   9.46   +0.03   +0.06   +0.06   +0.08   +0.10   +0.21   +0.33   10.00 −0.27 -  - −1.66 -  -  +0.33  \\x0c', '1404   L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411  After ablation,  the models surfaces and cross sections were analyzed by SEM-EDS to investigate the materials degradation.  3. Results and discussion  3.1. Microstructure and thermo-mechanical properties of the ceramics  The HfC-based sample (HCM) sintered at 1900 C reached the complete densification. The fractured and polished sections in Fig. 3a and b show a fine and homogeneous microstructure. The matrix mean grain  size was around 2.4  \\u242em. The MoSi2 phase was distributed among the matrix grains and formed pockets of 0.5  \\u242em. Additionally, scattered pockets of SiC/Si-C-O \\u242em were detected  based phases with dimension up  to 5  in  the matrix in amount around 2-3 vol%. In agreement with previous studies12,13 these phases are due to reaction of MoSi2 with C or CO species. The TaC-based sample (TCM), sintered at 1800 C, reached the complete densification. No porosity and a dense and fine microstructure with a mean grain size around 1  \\u242em was detected by SEM  on  the  fractured  and  polished  surface  in Fig.  3c and d. MoSi2 was distributed among  the grains and  formed pockets  of  3-4  \\u242em. As  for  the  previous  composite,  about 2-3 vol% of SiC/Si-C-O based phases were observed  in  the matrix. Basic mechanical properties12 and newly measured thermal properties  are  summarized  in  Table  3. Generally  speaking,  except  for  the  hardness,  higher  for HCM, TCM was more  thermo-mechanically  performing  than HCM,  possessing higher  toughness, strength and  thermal conductivity. The latter  increased  for  both materials  up  to  1500 C. The  lin−6 K −1 , ear CTE was  similar  for both  ceramics, 6.5-7   10 and  not  notably  different  from  other  transition  metal carbides.22           ×  3.2. Preliminary oxidation tests at 1600     C/5 min  Before the ablative tests on HCM and TCM, their oxidation behavior was analyzed. Mass variations were very different, i.e. +0.3% for HCM, +6.7% for TCM and  indicate  that TaC-based sample was more strongly oxidized  than HfC-based one. The higher oxidation  that TCM suffered  is also evident observing the damaged shape of the oxidized sample in the insets of Fig. 4 and the surface features in the SEM-images (Fig. 4a and b). According to XRD (Table 4), HfO2 was the only crystalline phase present on the surface of HCM. By SEM-EDS analyses, this sample was covered by HfO2 -based micro-cracked scale with scattered Mo-Si-O pockets. It can be observed  that 15% of MoSi2 is not enough to ensure a silica coverage of the HfO2 oxide. As for TCM, Ta2O5 crystals were detected by XRD. By SEM-EDS it was ascertained that Ta2O5 crystals are embedded in a silica-based glassy phase,  that formed  just on  the surface. The oxidation of this composite was dominated by formation of 1560 a  liquid phase due  to  the Ta2O5-SiO2 eutectic at  C.24 Presence of liquid phase at 1600 C enhanced dramatically the        T  a  b  l  e   3  T  h  e  r  m  o   m  c e  h  a  n  i  a c  l   p  r  p o  e  r  i t  e  s  :   V  i  c  k  e  r  s   m  i  c  r  h o  a  r  n d  e  s  s  (  V H  1  .  0  )  ,   Y  g  n u o  m  u d o  l  u  s   (  E  )  ,   C  N  B   r f  c a  t  u  r  e   t  n h g u o  e  s  s  (  K  I  c  )  ,   4   p  t   fl  e  u x  r  a  l   s  t  r  e  g n  t  h   a  t   r  o o  m   (  σ  R  T  )   a  d  n  a  t   0 0 2 1     C   (  σ  0 0 2 1  )  ,   t  h  e  r  m  a  l   e  p x  a  n  s  i  n  o  c  o  e  f  fi  c  i  e  n  t  (  λ  5 2  -  0 0 3 1  )  ,   t  h  e  r  m  a  l   c  u d n o  c  i t  v  t i  y   (  K  T  H  )  .  L  a  b  e  l   V H  1  .  0  (  G  P  a  )  E   (  G  P  a  )   K  I  c  (  M  P  a   m  1  /  2  )  σ  R  T  (  M  P  a  )   σ  0 0 2 1  (  M  P  a  )  λ  0 0 3 6 1 - 5 2  (  0 1  −  K  −  1  )  K  T  H  0 2  (  W  /  m   K  )  K  T  H  0 0 5  (  W  /  m   K  )  K  T  H  0 0 0 1  (  W  /  m   K  )  K  T  H  0 5 2 1  (  W  /  m   K  )  K  T  H  0 0 5 1  (  W  /  m   K  )  K  T  H  0 0 9 1  (  W  /  m   K  )  H  C  M   9 1  .  6   ±  ±   0  .  5  *  1  5 4  ±  ±   5  *  3  .  0  8  ±  ±   0  .  3 0  *  7  1 4  ±  ±   8 3  *  4  9 2  ±  ±   9 3  *  7  .  6  2  2 2  .  0   4 2  .  5   9 2  .  5   3 3  .  7   6 3  .  3   -  T  C  M   4 1  .  5   0  .  3  *  0  9 4  5  *  4  .  0  7  0  .  0  1  0  0 9  3 3  *  7  3 5  5 4  *  6  .  1  5  4 2  .  1   5 3  .  0   2 4  .  7   -   7 4  .  5   7 4  .  3  *  R  e  f  .   2 1  .  \\x0c', 'L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411   1405  Fig. 3. HCM (a, b) and TCM (c, d) fractured and polished sections.  Fig. 4. Surface features of HCM (a) and TCM (b) after oxidation test at 1600     C for 5 min. Macroscopic appearance of the samples is shown in the insets.  oxidation phenomena, eventually  leading  to bubble  formation and sample deformation, as shown in the picture of Fig. 4.  3.3. Ablation tests  Two  test campaigns with different flame  intensity were carried out. The  typical procedure  includes  the exposition of  the  sample at prefixed level of gas flow rate (in our test, stoichiometric compositions have been used) for 4 min. The main conditions and results are summarized in Table 2. As previously mentioned, graphite models were used  as reference materials. The  temperature profiles  for  the different samples during  the  test are shown  in Fig. 5. Temperature measurements were obtained  from  the  infrared camera  (field  Table 4  XRD and EDS results for HCM and TCM after oxidation and ablative tests: surface analysis.  Oxidation  type  HCM   TCM  Crystalline phases XRD   Chemical compounds  Crystalline phases XRD   Chemical compounds  Bottom up loading  furnace  Flame (test 1)   Flame (test 2)   43-1017:HfO2  43-1017:HfO2  43-1017:HfO2  SEM-EDS  HfO2 , SiO2 , Mo5 Si3 traces  HfO2 , HfCO, SiCO, SiO2  HfO2 , HfCO, SiCO, SiO2  25-0922:Ta2O5  21-1198:Ta2O5 25-0922:Ta2O5 21-1198:Ta2O5 25-0922:Ta2O5  SEM-EDS  SiO2 , Ta2O5 , (Ta, Si)O  SiO2 , Ta2O5 , (Ta, Si)O, Mo5 Si3 SiO2 , Ta2O5 , (Ta, Si)O, Mo5 Si3  \\x0c', '1406   L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411  Fig. 5. Temperature profile during   the ablative   tests: (a) oxygen flow rate = 0.035 g/s, fuel flow rate = 0.01 g/s, P = 500 W, (b) oxygen flow rate = 0.07 g/s, fuel flow  rate = 0.02 g/s, P = 1000 W.    ×                       of view/min focus distance: 24  18 /0.3 m, spatial resolution (IFOV): 1.3 mrad) pointed  toward  the  lateral  side of  the  test sample. Emissivity was evaluated  through a correlation with the pyrometric measurement. During  test 1,  the  temperatures increased  to 1327 C  for graphite  and  to 1527  and 1627 C for HfC and TaC,  respectively. On  the contrary during  test 2, graphite reached 1627 C, HfC 1927 C, whilst TaC was around 1727 C. Good  reproducibility of  the  tests was obtained  in  the case of condition 2 for TCM (see Fig. 5b and Table 2). The experiment shows similar thermal response when the same sample is exposed at equal “nominal” flame conditions (maximum temper50 ature difference  C). A very small mass  loss (0.23%) was detected during  the second exposition of  the specimen, probably due  to a spallation of  the oxide  layer. During  the first  test campaign, conversely to graphite, no appreciable mass variation was detected for the two ceramic materials (Table 2). During the second campaign, about 26% of graphite sample was  lost due to ablation. Conversely, the ceramic models only experienced a modest weight increase (+0.1-0.3%), see Table 2. The recorded values of  temperature depend on  the radiative properties of  the samples,  i.e. on  their emissivity, and  in  turn this property may change during  temperature  increase, due  to changes in the microstructural features. The low value of the maximum temperature obtained for the graphite is due to the high emissivity of the material, 0.85. For hafnium carbide and  tantalum carbide, emissivity values measured  in vacuum are  lower  than  those of graphite, about 0.6 respectively.23 During  and 0.3 at 1200 C,  the first campaign, the emissivity of hafnium carbide was similar  to  that reported in Ref.23 , while a notable increase was observed for TaC. During the second campaign, TaC emissivity almost reached that of graphite whilst HfC emissivity was still around 0.65. As a result of the higher emissivity, the maximum temperature reached on the TaC surface was lower than for HfC. Fig. 6 shows the samples after the ablative tests. At glance, it can be observed that the graphite-based samples suffered a high degradation after the exposure to the ablative flame. In particular,     after  the second  test  the dimensions of  the graphite sample are visibly reduced compared  to  those of  the sample after  the first test. For both ceramic models,  the effect of high  temperature exposition was only a  slight oxidation. Visually,  the ceramic samples changed their color from gray (for HfC) or brown (for TaC) to white, due to formation of an oxide layer. Very likely the formation of an oxidized layer was responsible for the increase of TaC emissivity, which in turn, determined a peak temperature, lower than for HfC and close to that of graphite (Fig. 5).  3.4. Microstructural characterization after ablation tests  Fig. 7  shows  the  surface  features of  the graphite  samples before and after  the  tests 1 and 2.  It  is apparent  that a progressive degradation of graphite-based materials occurred  in agreement with data reported in Table 2 and photos of Fig. 6. At the microstructural level, surface erosion occurring during high temperature tests caused a dramatic change of the microstructure from dense to increasingly porous. According to X-ray diffraction analyses (Table 3), for HCM, the oxide  layer was based on HfO2 whilst for TCM  the oxide layer was based on Ta2O5 . The SEM-images in Fig. 8 show the microstructural features of the surfaces of HCM (a and b) and TCM (d-g) after ablation tests and  the corresponding EDS spectra (c,  insets  in d and f). The most evident effect of high temperature exposure was oxidation despite the presence of butane/propane in the flame. After both tests the hafnium oxide layer on the surface was not continuous and Si-C-O regions were occasionally found in proximity of cracks/voids  (Fig. 8a and b). HCM sample exposed  to  test 2 showed a slight grain coarsening due  to exposure  to higher temperature (Fig. 8b). The surface of TCM after  test 1  is displayed  in Fig. 8d and looks homogeneously composed by Ta2O5 grains with glassy Ta-Si-O phases and cracks. When TCM was  tested  in more severe conditions (test 2), the external surface showed areas with dense  tantalum oxide and areas where ablation occurred  leaving silica exposed  (Fig. 8e). Tantalum oxide showed  lamellar  \\x0c', 'L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411   1407  Fig. 6. Samples after the ablative tests 1 and 2: frontal view on the left and side view on the right.     voids, probably caused by  the  lifting of  the platelet grains by gas developed underneath  the outermost  layer. In addition,  the dense morphology of tantalum oxide in Fig. 8f is due to the melting of Ta2O5-SiO2 phase, liquid above 1560 C.24 On the other side, the glassy phase showed that nano-aggregates of tantalum oxide precipitate and  tend  to coalesce  leaving  the silica-based glass darker,  i.e. progressively  less viscous, Fig. 8g. It  is worthy  to note  that on  the surface of both materials no continuous silica-based layer was observed, as explained later. As for the cross sections of HCM and TCM, formation of an oxide scale is well evident. Pictures in Fig. 9 show the sections of  the samples after  test 2. Oxidized  layer  thickness was not constant through the sample profile. The reason for the thickness variability resides in the ablation which was more pronounced on the frontal region than on the lateral regions. This also suggests that weight data reported in Table 2 for the carbides are the result of simultaneous weight gain due  to oxidation and weight  loss due  to oxide ablation. Owing  to  the  removal of oxide on  the top of the models, comparisons on the oxidation degree will be based on  the  lateral oxide  thickness. More  in detail,  the  lateral thickness of the oxide scale was comparable after tests 1 and 2 for HCM, 0.06-0.10 mm, whilst for TCM, it increased from 0.06  to 0.33 mm after test 2 (Table 2). These tests confirmed the lower oxidation rate of HfC compared to TaC, which is consistent with results of conventional oxidation  tests  reported  in Section 3.2 and with previous works.12 In  the central areas, HCM showed three main distinct  layers:  the original bulk, a Hf-C-O scale with SiC and Si-O-based phases and an outermost HfO2 layer composed by a porous inner one and a dense outer one, Fig. 9b. As for TCM, Fig. 9d, only a thick Ta2O5 layer was composing the external scale, which was, in addition, not well adherent to the unreacted bulk. Fig. 10 compares  the magnified section of HCM specimens after  tests 1 (Fig. 10 a) and 2 (Fig. 10b-e)  taken  in  the frontal region. Some damage of the scale was caused by polishing but also  indicates  that  the oxide was quite brittle. It can be easily seen that in the central area analyzed the total thickness passed from 55  \\u242em after exposure at 1527 C to around 70  \\u242em after the second test at 1927 C and the oxide scale started to be not well adherent only  in  the harsher conditions.  In addition after  test 1, Fig. 10a,  the modified  layer consisted of dense HfO2 grains contoured by Hf-C-O spurious phase, SiC grains and Si-C-O glassy phases spread throughout the microstructure. Only after test 2 a thick continuous Hf-C-O layer formed, Fig. 10b, which        Fig. 7. Graphite surface features before (a) and after ablation test 1 (b) and 2 (c).  \\x0c', '1408   L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411  Fig. 8. HCM surface features after test 1 (a) and test 2 (b), TCM surface features after test 1 (d) and test 2 (e-g) and the corresponding EDS spectra.  was well adherent  to  the unoxidized bulk  (Fig. 10d) but not firmly anchored  to  the upper porous HfO2 scale. It  is also visible  that  the oxidation started from HfC grain boundaries and progressively attacked  the carbide grains, Fig. 10e. Moreover, oxidation of HfC was accompanied by formation of carbon, visible as a dark contrast phase  in Fig. 10f. The good anchorage between HfC and Hf-C-O was probably due  to similar coefficient of thermal expansion and crystal structure, but also to the nano-sized morphology of  the mixed phase which could precisely follow the ceramic profile. In this Hf-C-O scale, MoSi2 oxidation products could still be found, as the EDS spectrum and Fig. 10d confirm, however, due to the beam limitations, we cannot ultimately say if this phase was pure molybdenum, derived from MoSi2 dissociation, or molybdenum oxide. Moving further to the top, hafnia grains started to coarsen, coalesce and form a denser scale which prevented fast ceramic oxidation, Fig. 10c.        SEM images of the cross section of TaC-based ceramics are shown in Fig. 11. For this composite, the modified thickness in the central zone passed from 77  \\u242em after exposure at 1627 C to 130  \\u242em after exposure at 100 C higher, which  testified a notably  faster oxidation process as compared  to HCM. After test 1  the oxide scale was mostly crack free. In Fig. 11a, starting from  the bulk, a dense Ta2O5 layer with MoSi2 oxidation products was  topped by a  tiny porous Ta2O5 scale. Silica was recognizable as dark phase scattered in the whole modified scale, whilst MoSi2 disappeared at 30  \\u242em from  the surface. Close  to the bulk, Mo5Si3 phase was  recognizable around MoSi2 with brighter contrast,  see Fig. 11c and EDS  spectra. After exposure to test 2, the modified scale basically maintained the same features as  those described  for  test 1, but  just grew  in  thickness, Fig. 11b. In addition  to  the previous  test, a dense Ta2O5 layer formed on the outermost surface, Fig. 11d, and SiO2 and  Fig. 9. Pictures of the cross sections of the ceramics after ablation test 2, HCM (a, b) and TCM (c, d). (a, c) Curving zone of the samples, and (b, d) frontal zone.  \\x0c', 'L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411   1409  Fig. 10. HCM cross sections after test 1 (a) and test 2 (b) with microstructural details (c-f) and EDS spectra.  Mo-Si phases progressively disappeared moving from the bulk outwards, Fig. 11e-g.  3.5. Oxidation mechanisms during ablation tests  The main combustion products developed during the ablation tests are CO/CO2 (75%) and H2O (25%). Although the oxygen partial pressure  is negligible,  the final effect on  the carbides microstructure  is oxidation. Thus we must conclude  that mass loss recorded for some  tests or reduction of  the sample height (see Table 2) is due to erosion of the scale, because pure oxidation (see Section 3.2 and Refs. 12,25) would  lead  to a net weight increase.  As  for  the  oxidation  of  pure  hafnium  carbide,  several studies25-27 indicate that oxide growth of carbides is parabolic indicating protective oxidation behavior. During  the oxidation, gaseous CO is produced as a by-product and introduces porosity which allows diffusion via pores. Gas formed below  the oxide layer can also lift and disrupt the oxide layer. Nevertheless, the generated oxide  layer  is generally considered  to be partially protective.25-27 Bargeron et al.26 reported the formation of a layered oxide, comprising a porous outer HfO2 layer and a dense HfO2−xCy , interlayer that represents the real diffusion barrier for the oxygen. The effect of MoSi2 addition  to HfC was already reported  in several studies8,12,25 and  it was shown  that  in  the outer layer, silica formed by oxidation of MoSi2 partially filled  \\x0c', '1410   L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411  Fig. 11. TCM cross sections after test 1 (a) and test 2 (b) and microstructural details of the sections (c-f).  the pores, making  the scale more compact, whilst  in  the  intermediate HfO2−xCy layer, SiOxCy and MoO2 species  formed. Noteworthy,  in none of  the cases analyzed, a continuous silica layer was observed on the surface of HfC samples. For the present experiment, the situation has several similarities with conventional oxidation: after test 1, no clear formation of an intermediate oxycarbide layer is observed, but dense HfO2 grains are contoured by Hf-C-O spurious phase, SiC grains and Si-C-O glassy phase. After test 2, a thick continuous Hf-C-O interlayer is formed. In this layer, we also detected Si-O-C and Mo-O phases deriving  from oxidation of MoSi2 .8 However, after  test 2  the outer  layer was more porous even despite  the presence of silica or SiC-based phases. No silica was found on the surface. Several can be  the reasons for such behavior -  the amount of SiO2 formed  is not enough  to fill  the voids  in  the        oxide layer and cover the oxide surface - silica is formed but is actively oxidized to SiO(g) by reaction with water vapor to form a volatile Si(OH)4 , as typical of combustion environment.28 As  for TaC, oxidation experiments  in previous works12 at 1650 C and  this work at 1600 C revealed quite dramatic oxidation phenomena:  (1)  large volume expansion associated  to the oxidation of TaC  to Ta2O5 ;  (2) partial detachment of  the oxide scale from  the unreacted bulk due  to  thermal expansion coefficient mismatch; (3) fast consumption of the ceramic. The addition of MoSi2 did not improve the oxidation resistance due 1560 to  the Ta2O5 -SiO2 eutectic at  C  that caused presence of boiling liquid phase on the sample surface. Ablation  tests  led  to  less dramatic oxidation mechanisms, despite the temperatures were higher than 1560 C. The oxidation products were  those expected  from oxidation of  the bulk        \\x0c', 'L. Pienti et al. / Journal of the European Ceramic Society 35 (2015) 1401-1411   1411  constituents,  e.g. Ta2O3 from TaC, SiO2 and Mo5Si3 from MoSi2 . Again, silica pockets were  found  to fill  the porosities of the scale but no continuous layer was found on the surface, as for HfC. Sporadic pockets of silica with small precipitated crystals of Ta2O5 were found on the surface but the sample damage was much less pronounced compared to conventional oxidation. In conclusion, whatever is the leading phenomenon for these carbides, replacing bulk compact material with a partially porous oxidized  layer makes  the material more prone  to erosion during exposition at high energy fluxes  such as  those occurring in propulsion conditions,  thus oxidation phenomena should be avoided as much as possible. However,  it  should be  recalled that propulsive  conditions  are  even more  reducing  than  the present ones,  thus deleterious effects of oxidation  should be more limited.  4. Conclusions  HfCand TaC-based composite materials containing 15 vol% MoSi2 as sintering aid were selected to produce prototypes for ablation tests in a mixture of oxygen/butane/propane. The ablation resistance of the HfCand TaC-ceramics was compared to that of graphite. SEM analysis of surfaces and sections was conducted to study the material degradation. The composites have a superior ablation resistance compared to graphite, showing only slight surface oxidation and weight variation  for  temperature above 1700 C. The graphite components underwent strong erosion, with reduction of the initial dimensions. On the other side, carbide components did not appreciably change their dimensions compared  to  the pristine ones. Although TaC based composite has superior  thermo-mechanical properties,  it  is more prone  to oxidation than HfC. However, the formation of tantalum oxide resulted in an increased emissivity which allowed the heat to be re-radiated away faster, similar to the behaviour of graphite.     Acknowledgements  D. Dalle  Fabbriche  and A.  Piancastelli  acknowledged for their technical support.  are   gratefully  References  1. Courtright EL, Graham HC, Katz AP, Kerans RJ. Ultra high   tempera ture   assessment   study:   ceramic matrix   composites.   In: WL-TR-91-4061  (ADA262740). 1992.  2.   Johnston JR, Signorelli RA, Freche JC, NASA technical note/D: 3428 Per formances of rocket nozzle material with several solid propellants. National  Aeronautics and Space Administration; 1966.  3. Hickman   R, McKechnie   T,   Agarwal   A.   Net   shape   fabrication   of  high   temperature materials   for   rocket   engine   components.   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Microstruc ture and properties of HfC and TaC-based ceramics obtained by ultrafine powder. J Eur Ceram Soc 2011;13:619-27.  15. Silvestroni L, Sciti D. Sintering behavior, microstructure, and mechani cal properties: a comparison among pressureless sintered ultra-refractory carbides. Adv Mater Sci Eng 2010;2010:1-11. Article ID 835018.  16. Khaleghi E, Lin Y, Meyers MA, Olevsky EA. Spark plasma sintering of tantalum carbide. Scr Mater 2010;63:577-80.  17. Ghaffari SA, Faghihi-Sani MA, Golestani-Fard F, Mandal H. Spark plasma  sintering of TaC-HfC UHTC via disilicides sintering aids. J Eur Ceram Soc 2013;33:1479-84.  18. Liu H, Liu L, Ye F, Zhang Z, Zhoua Y. Microstructure and mechanical prop erties of the spark plasma sintered TaC/SiC composites: effects of sintering temperatures. J Eur Ceram Soc 2012;32:3617-25.  19. Liu   JX,   Huang   X,   Zhang   GJ.   Pressureless   sintering   of   hafnium 2013;96:  carbide-silicon   carbide   ceramics.   J   Am   Ceram   Soc   1751-6.  20. Balani K, Gabriela G, Agarwal A, Hickman R, O’Dell   JS. Synthesis,  microstructural   characterization   and mechanical   property   evaluation  of 2006;89:  vacuum   plasma   sprayed   tantalum   carbide.   J Am Ceram   Soc   1419-25.  21. Pienti L, Silvestroni L, Melandri C, Landi E, Sciti D. Microstructure,  mechanical properties and oxidation behavior of TaCand HfC-based materials containing short SiC fiber. Ceram Int 2015;41:1367-77.  22.   Jun CK, Shaffer PTB. Thermal expansion of niobium carbide, hafnium car bide and  tantalum carbide at high  1971;24:323-7.  temperatures. J Less Common Metals  23. Sani E, Mercatelli L, Fontani D, Sans JL, Sciti D. Hafnium and tantalum carbides for high temperature solar receivers. J Renew Sustain Energy 2013;33,  0.33104-1/13.  24. Roth RS, Waring JL. Diagram 4448: Ta2O5 -SiO2 . In: Phase diagrams for  ceramists. 1970.  25. Charpentier L, Balat-Pichelin M, Sciti D, Silvestroni L. High   temperature  oxidation of Zrand Hf-carbides: influence of matrix and sintering additive. J Eur Ceram Soc 2013;33:2867-78.  26. Bargeron BB, Benson RC, Jette AN, Phillips TE. Oxidation of hafnium  carbide  in  the  temperature  1993;76:1040-6.  range 1400     C   to 2060     C. J Am Ceram Soc  27. Courtright EL, Prater JT, Holcomb GR, St Pierre GR, Rapp RA. Oxidation  of hafnium carbide and hafnium carbide with additions of  praseodymium. Oxid Metals 1991;36:423-37.  tantalum and  28. Opila JE. Oxidation and volatilization of silica formers in water vapor. J Am Ceram Soc 2003;86:1238-48.  \\x0c']"
},{
  "_id": 15,
  "PDF": "Ablation-resistant carbide Zr0.8Ti0.2C0.74B0.26 for oxidizing environments up to 3,000 degrees C.pdf",
  "Text": "['A RT IC L E  Received 3 Aug 2016 | Accepted 1 May 2017 | Published 14 Jun 2017  DOI: 10.1038/ncomms15836  OPEN  Ablation-resistant carbide Zr0.8Ti0.2C0.74B0.26 for oxidizing environments up to 3,000 °C  Yi Zeng1,2, Dini Wang1, Xiang Xiong1, Xun Zhang2, Philip J. Withers2, Wei Sun1, Matthew Smith2,  Mingwen Bai2 & Ping Xiao2  Ultra-high temperature ceramics are desirable for applications  in the hypersonic  vehicle,  rockets, re-entry spacecraft and defence sectors, but  few materials can currently satisfy the  associated  high  temperature  ablation  requirements. Here we  design  and  fabricate  a  carbide  (Zr0.8Ti0.2C0.74B0.26)  coating  by  reactive melt  infiltration  and  pack  cementation  onto  a C/C composite.  It  displays  superior  ablation  resistance  at  temperatures  from  2,000-3,000 °C, compared to existing ultra-high temperature ceramics (for example, a rate  of material  loss over 12 times better than conventional zirconium carbide at 2,500 °C). The  carbide  is  a  substitutional  solid  solution  of Zr-Ti  containing  carbon  vacancies  that  are  randomly occupied by boron atoms. The sealing ability of the ceramic’s oxides, slow oxygen  diffusion and a dense and gradient distribution of ceramic result  in much slower  loss of  protective oxide layers formed during ablation than other ceramic systems,  leading to the  superior ablation resistance.  1 State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China. 2 School of Materials, University of Manchester,  Manchester M13 9PL, UK. Correspondence and requests for materials should be addressed to X.X. (email: Xiong228@sina.com) or  to P.X.  (email: P.xiao@manchester.ac.uk).  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications  1  \\x0c', 'A RT IC L E  NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  F uture hypersonic aerospace vehicles offer the potential of a step jump in transit speeds. Currently, one of the biggest challenges is how to protect critical components such as leading edges, combustors and nose tips so that they survive the severe oxidation and extreme scouring of heat fluxes at temperatures in excess of 2,000 °C during flight1,2. The diborides of Hf and Zr are considered to be the most promising candidates for such components3,4, offering the best oxidation resistance up to 1,500 °C among candidate ultra-high temperature ceramics (UHTCs)5. In particular, ZrB2 has attracted much attention due to its low density and cost6-8. However, there are two critical factors hindering its application: first, a high level of boron (about 66 at. %) leads to severe loss of material under the scouring of hot gas because of the rapid evaporation of boron oxides at temperatures above 1,200 °C (refs 9,10), second, monolithic ZrB2 tends to fail catastrophically due to a combination of low toughness and poor thermal shock resistance11. To reduce the evaporation of boron oxides and to improve ZrB2 ablation resistance, much attention has been focused on adding silicides (for example, SiC12,13, MoSi2 (ref. 14) and so on) (for example, ZrC15) and carbides to ZrB2 to form multi-phase ceramics. By contrast relatively little attention has been directed towards the development of a single-phase ceramic comprising multiple elements. In particular, quaternary carbides with low boron contents for ablation resistance have not been reported since UHTCs were first proposed in 1930s16. Moreover, although the oxide of ZrC evaporates less at higher temperature due to the absence of boron, it is generally believed that monolithic zirconium carbide (for example, ZrC) has inferior oxidation resistance compared to the diborides (for example, ZrB2) making it a poor option for anti-ablation applications17-19. All of the above factors mean that the current number of candidate UHTCs for use in extreme environments is limited and it is worthwhile to explore the potential of new single-phase ceramics in terms of reduced evaporation and better oxidation resistance. In addition, it has been shown that introducing such ceramics into carbonfibre-reinforced carbon matrix (C/C) composites may be an improving thermal-shock resistance20,21. effective way of Here a coating of the quaternary carbide, Zr0.8-Ti0.2-C0.74-B0.26, laid down on a C/C composite by reactive melt infiltration (RMI) and pack cementation (PC) is proposed (Methods section and Supplementary Fig.1). The carbide comprises a substitutional solid solution with low boron content. In addition, to improve the thermal-shock resistance, and to decrease the risk of cracking of the carbide coating during ablation, some carbides are allowed to infiltrate into the C/C composite. The experimental results presented here suggest the carbide coating displays better ablation resistance at 2,000-3,000 °C than existing candidate UHTCs such as Zr-based carbide and diborides and other high temperature composites. More broadly, this work provides a platform for building a series of UHTCs based around the group IV/V transition metals.  Results  performance.  Figure  Zr0.8Ti0.2C0.74B0.26  A profile of ablation 1 compares the ablation resistance of coating on C/C composite alongside other common UHTCs and composites. The mass ablation rate (MAR) and linear ablation rate (LAR) characterize the mass loss and dimensional stability of the materials, respectively. In general, a high MAR (that is, rapid loss of mass) and LAR (that is, rapid degradation of surface integrity) indicate poor ablation performance. Hence, the MAR and LAR results for our carbide in Fig. 1 demonstrate a significant improvement in ablation resistance relative to existing UHTCs coatings or composites as well as monolithic ZrB2-SiC ceramics  39.2  28.0  16.8  5.6  0.0  -2.3  -4.6  )  1 -  s  m  μ  (  R  A  L  C/C-Zr0.8Ti0.2C0.74B0.26,3,000°C,60s C/C-Zr0.8Ti0.2C0.74B0.26,2,500°C,60s C/C-Zr0.8Ti0.2C0.74B0.26,2,000°C,120s C/C-Zr0.8Ti0.2C0.74B0.26,2,000°C,60s C/C-Zr0.8Ti0.2C0.74B0.26,2,000°C,120s ZrB2-SiC,3,000°C,100s C/C-ZrTiC/SiC,3,000°C,60s C/C-Zr0.83Ti0.17C,3,000°C,60s C/C-Zr0.83Ti0.17C,2,500°C,60s C/C-Zr0.83Ti0.17C,2,000°C,60s C/C-Zr0.57Ti0.43C,3,000°C,60s C/C-Zr0.57Ti0.43C,2,500°C,60s C/C-Zr0.57Ti0.43C,2,000°C,60s C/C-MoSi2-SiC-TiSi2. 2,471°C,300s C/C-ZrC,2,500°C,60s C/C-3,000°C,60s C/C-2,500°C,60s C/C-SiC-TaC,2,000°C,60s C/C-SiC-2,000°C,60s  0.0  0.5  1.0  1.5  MAR (mg cm-2 s-1)  Figure 1 | Ablation performance of Zr0.8Ti0.2C0.74B0.26 composites.  Comparison of the ablation rates (MAR and LAR) for a range of candidate UHTC composites including: ZrC (RMI)29, SiC (RMI), SiC-TaC (CVI)42, Zr0.83Ti0.17C (RMI)29, Zr0.57Ti0.43C (RMI)29, ZrTiC-SiC (RMI), MoSi2-SiCTiSi2 (RMI)43, ZrB2-SiC (SPS) and C/C composites (CVI), as well as the  Zr0.8Ti0.2C0.74B0.26 studied here (fabrication methods shown in brackets).  Above ablation tests were conducted by authors in oxyacetylene machine.  fabricated by spark plasma sintering (SPS). For instance, the LAR of ZrC at 2,500 °C is 8.0 mm s \\x00 1 and MAR is 1.10 mg cm \\x00 2 s \\x00 1, 3.5 mm s \\x00 1 whereas our carbide gains in thickness and 0.14 mg cm \\x00 2 s \\x00 1 in weight. This is because the oxide layer expands and increases the weight countering any material loss from ablation, indicating an almost negligible loss of our carbide. This is over 12 times better than the loss of material for ZrC (assuming both carbides have the same volume expansion of oxides per second). It is noteworthy that the MARs of our carbide 2,000-2,500 °C distribute from 60-120 s at around the zero, indicating a slight weigh loss or weigh gain. These higher negative values of LARs obviously distinguish it from other UHTCs indicated by the green arrows in Fig. 1, indicating that a good quality protective oxide layer of carbide is formed during the ablation experiment (Protective mechanisms section). This layer is strongly adhered to the C/C composite substrate being able to endure the scouring of the high speed hot gas and providing a high level of protection to the substrate (Fig. 3). At 3,000 °C, our carbide still exhibits low LAR and MAR. In addition, it should be noted here that the ablation performance can be mainly attributed to the ceramics, regarding the RMI/chemical vapour infiltration (CVI) process (Methods section and Supplementary Fig.1 and Supplementary Note 1). Generally, a ceramic coating would remain on the composites fabricated by the above methods. Once the coating has been depleted, the carbon matrix would be exposed and be detached quickly by hot gas, causing a very high LAR (see the significance of change for LAR of the composites in Fig. 1). However, the MAR change would be less because of the weight gains of oxides from the ceramics. Figure 2 shows a photograph of the ablation test and the morphology of the tested sample. Despite the low level of boron present, our carbide displayed the characteristic light green flame during the ablation test, which in contrast to the carbide (that is, Zr0.8Ti0.2C fabricated by RMI), and typical of the ablation flame of ZrB2 (ref. 22), as shown in Fig. 2a. Generally, borides (for example, ZrB2) show better levels of oxidation resistance than ZrC)23. their carbides (for example, From the comparison between Zr0.8Ti0.2C0.74B0.26 and Zr0.8Ti0.2C shown in Fig. 1, it is inferred that the improvement in ablation resistance can be  2  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications      \\x0c', 'NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  ART IC L E  a  b  Before test  2,000 °C, 120s  2,500 °C, 120s  Test of Zr0.8Ti0.2C0.74B0.26  c  5.0 4.5 4.0 3.5 3.0 2.5 2.0 1.5 1.0  0.5  0.0 0.0  1.5  1.0 0.5  2.0 2.5  Test of Zr0.8Ti0.2C  6.0 4.55.05.5  3.0 3.5 4.0  m  m  3,000 °C, 60s  72 μm 60 50 40 30 20 10 0 -10 -20 -30 -40 -50 -60  -76 μm  Figure 2 | Photograph of ablation test and morphology of tested sample. (a) Green and orange flames are seen in tests of Zr0.8Ti0.2C0.74B0.26 and  Zr0.8Ti0.2C, respectively. (b) Comparison of surface of the 30 mm diameter samples before and after ablation. Black-gray sample is before test, and middle and right samples experienced 120 s ablation of 2,000 °C and 2,500 °C, respectively. (c) Surface profile of central region of sample ablated at 3,000 °C, showing the ablated traces (some convexities with the rises of o72 mm) due to evaporations of oxides with low melting points (see Protective mechanisms). But no ablated hollows appeared on surface.  a  1.0  0.8  0.6  0.4  0.2  0.0  b  )  %  (  n o  i  t  i  s  o p  m  o  c  e  s  a h  P  c  Pore  Carbon  Ceramic  1,000  3,000 5,000 7,000 9,000 Distance from top surface (μm)  11,000  Transitional region  Top region  Figure 3 | Morphology of C/C-Zr0.8Ti0.2C0.74B0.26/SiC via X-ray computed-tomography. (a) Overall view of composites showing carbides (yellow), carbon fibres and pyrocarbon (grey) and pores (red). Due to the resolution limit, only pores larger than 1,000 mm3 are quantified. (b) Distributions of the  phases from the top to the bottom of the sample determined from the X-ray computed tomography (CT) virtual slices. The dots are the volume percentage  of each phase calculated from each virtual slice. The solid lines are the fitted curves. The zig-zag shape of the dot is due to the alternating plies. The grey  area represents the top of  the sample which has been excluded from quantification due to strong artefacts. (c) Virtual CT cross-sections showing the  transitional region comprising the carbon matrix and ceramic (40 vol. %) and a region near the top (ceramic: 76 vol. %). Scale bar, 5 mm.  mainly attributed to the introduction of boron into Zr0.8Ti0.2C. In addition, as shown in Fig. 2b,c, the ablated surfaces are relatively smooth and free from any obvious erosion hollows and cracks. The protective oxide layer grows with the increasing temperature remaining essentially intact throughout. Consequently, the Zr0.8Ti0.2C0.74B0.26 carbide exhibits a level of ablation resistance and protection not seen in other Zr-based carbides and diborides and common high-temperature composites. The results also  suggest that C/C composite modified by (Figs 2 and 3) displays good thermal-shock resistance.  Zr0.8Ti0.2C0.74B0.26  the Zr0.8Ti0.2C0.74B0.26  Microstructure and constituents. Figure 3 shows the distribution of ceramic, the carbon (carbon fibre and pyrocarbon), and the pores below the surface into the test-piece. The surface region comprises up to 75% ceramic and 25% C. It is  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications  3      \\x0c', 'A RT IC L E  a  Scanning line   NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  Top layers  Ceramic/carbon composites  b  )  %  . t  a  (  n o  i  t  c  a  r  F  60  50  40  30  20  10  0  100  80  60  40  20  0  0  500  1,000 1,500 Distance (μm)  2,000  2,500  0  500  1,000 1,500 Distance (μm)  2,000  2,500  Zr  Ti  Si  C  B  Figure 4 | Microstructure and element distribution of carbides. (a) SEM image of cross-section of top ceramic layers and ceramic/carbon composites.  From left to right, the external  layer is the carbide with the thickness of 100-200 mm, the second layer composed of the carbide and SiC with the thickness  of 200-300 mm. Beneath that there are the ceramic and carbon composites as indicated by the arrows. (b) Electron probe microanalysis reveals element  distributions. Lines of pink, blue, red, purple and olive lines represent Zr, Ti, Si, C and B, respectively. In the ceramic areas (white areas in a), the ratios of Zr,  Ti, C and B are about 40, 10, 37 and 13 at. %, respectively. Grey areas beneath the external  layer are SiC according to results of the scan. Scale bar, 500 mm.  sufficiently dense that the top surface (see cross-section in Fig. 4) acts as a barrier to resist oxidation and scouring from the hot gas during the ablation test. The porosity (pore size 410 mm) is o3% and the volume of ceramic is higher than 58 % near the free surface shown in Fig. 3b,c (top region). Deeper into the sample (7,000 mm), it comprises carbon and pores (shown in Fig. 3a,b). In these areas (carbon-based part), the volumes of carbon and pores reach 90 and 10%, respectively, with no ceramic present. However, in the transitional region shown in Fig. 3c, the ceramic and carbon show a smooth gradient distribution with no sharp interface between the ceramic and the C/C composite. Generally, the C/C composite and Zr-based ceramics have significantly different coefficient of thermal example, CTEC/C ¼ 0.38-2.18 \\x02 10 \\x00 6 C \\x00 1 expansion (CTE) (for 24), CTEZrC ¼ 6.70 \\x02 10 \\x00 6 C \\x00 1 and CTEZrB2 ¼ (ref. (ref. 25) 6.66-6.93 \\x02 10 \\x00 6 C \\x00 1 (ref. 26)). Thus the gradient distribution employed in this work, together with the fibre reinforcement and weakened pyrocarbon interfaces27,28 might alleviate the mismatch in CTE, leading to an improvement of thermal shock resistance and the density of whole sample because of the presence of the carbonbased part. Figure 4 shows a cross-section through the top layer of carbide. The ablation resistant layer composed is about 100-200 mm in thickness. Beneath the carbide layer, SiC is identified which has formed from the reaction between the original carbon and the Si/SiO during PC process (Methods section and Supplementary Note 2), as shown in Fig. 4a. The formula of carbide was obtained, according to elemental analysis conducted by electron probe microanalysis (EPMA) as shown in Fig. 4b. The ratio of Zr and Ti is dictated by the raw powders (80 at. %:20 at. %) which was optimized for ablation resistance at temperatures over 2,000 °C from our previous investigation on the doping effect of Ti in Zr29. The ratio of B to C was 0.74:0.26 according to the EPMA results which is discussed in detail in the following section.  of Zr0.8Ti0.2C0.74B0.26  Zr0.8Ti0.2C0.74B0.26  Zr0.8Ti0.2C(1 \\x00 x)  Zr0.8Ti0.2C(1-x)  Zr0.8Ti0.2C0.74B0.26  Actually, the formation of our carbide forms in three stages. First, the Zr-Ti melt infiltrates the porous C/C composite at high temperature and reacts with the pre-deposited pyrocarbon30 to (0oxo1) form (Methods section and Supplementary Fig. 2). Subsequently, the forms through the boration of Zr0.8Ti0.2C(1-x) via solid diffusion of boron atoms (see following reactions) during PC process. is a substitutional solid solution which has a FCC structure along with carbon vacancies (Fig. 5b,e). In this process, the boron atoms fill the original vacancies replacing the carbon by diffusion, which has not changed the stacking type of atoms of Zr-Ti carbide. This interpretation is substantiated by ZrC diffraction peaks, instead of ZrB2, obtained from the top surface and cross-section of the composites shown in Fig. 5a, which has been also confirmed by the high resolution TEM images and their diffraction patterns of FCC shown in Fig. 5e,f. However, the replacement of interstitial atoms causes a variation in the crystal lattice constant (a) shown in Fig. 5b,e and f. For instance, the absence of carbon atom and substitution of Zr by Ti in Zr0.8Ti0.2C(1 \\x00 x) result in a smaller a, as shown in Fig. 5b, compared with pure ZrC. Nevertheless, the vacancies were occupied by the boron atoms leading to a slight increase of a. Consequently, the structural changes in Zr-Ti-C-B carbide can be described using the schematic representation shown in Fig. 5c,d. The defect channels in the FCC crystals of Zr0.8Ti0.2C(1 \\x00 x) have been built by disordered carbon vacancies (possibly having short-range-order31, Fig. 5c). The small boron atoms diffuse relatively quickly through the crystal boundaries, interstices and the defect channels to preferentially occupy the vacancies originally occupied by carbon atoms (Fig. 5d), due to the lower formation energy of a carbon vacancy31. This suggests that the large extent of the boron distribution in the carbide shown in Fig. 4 is possibly attributed to the pre-existing crystal defects that greatly promoted the diffusion of boron atoms. It is  4  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications    \\x0c', 'also inferred that, due to the pre-synthesized stable FCC structure during the RMI, the has been inhibited to transform into hexagonal structure of crystals such as ZrB2 and TiB2 during the subsequent reactions. Consequently, according to the variations of structure and phases as well as the raw materials, we conclude that the following boration reactions occurred during the PC process.  Zr0.8Ti0.2C(1 \\x00 x)  in Zr0.8Ti0.2C0.74B0.26,  14Zr0:8ð1 \\x00 xÞTi0:2ð1 \\x00 xÞCð1 \\x00 xÞ  \\x01 x V0  CZr0:8Ti0:2ðsÞ þ 3xB4CðsÞ þ xB2O3ðl:gÞ ¼ \\x01 x Zr0:8Ti0 :2 \\x01 B 14Zr0:8ð1 \\x00 xÞTi0:2ð1 \\x00 xÞCð1 \\x00 xÞ 14Zr0 :8Ti0 :2Cð1 \\x00 xÞBxðsÞ þ 3xCOðgÞ :  ð  ÞðsÞ þ 3xCOðgÞ ¼  ð1Þ  7Ti Zr  ð  Þðs:lÞ þ 3B4CðsÞ þ B2O3ðl:gÞ ¼ 7Zr Ti ÞB2ðsÞ þ 3COðgÞ :  ð  ð2Þ  7CðsÞ þ 2B2O3ðl:gÞ ¼ B4CðsÞ þ 6COðgÞ : where V 0 x \\x01 Zr0.8Ti0.2 and are the vacancy of carbon atoms and extra metal atoms in non-stoichiometric carbide, respectively. In reaction (1), Zr0:8ð1 \\x00 xÞTi0:2ð1 \\x00 xÞCð1 \\x00 xÞ is equivalent to Zr0.8Ti0.2C(1 \\x00 x) formed from the RMI process. B4C and B2O3 are the raw materials and Zr0.8Ti0.2C(1 \\x00 x)Bx is the final  ð3Þ  C  \\x01 xV 0  C \\x01 xZr0:8Ti0:2 ;  product. In this work, the average content of boron is about 13 at. % in the carbide and thereby x is 0.26 here. Hence, reaction (1) governs the main reaction of boron atoms. V 0 C finally disappeared because of the occupation of boron atoms. Moreover, the residual metal such as Zr32 and Ti in the composite would further react with the raw materials to form diborides (see the diborides observed by EPMA in Supplementary Fig. 4), as shown in reaction (2). Actually, the doped B2O3 may react with carbon fibres and pyrocarbon, as shown in reaction (3), which is another source of the B4C in reaction (1). Reaction thermodynamics are referenced in Supplementary Fig. 5.  Protective mechanisms. Figure 6 shows the surface and crosssectional microstructure at the centre of the ablated surface and the phases across the whole ablated surface. At 2,000 °C, and SiC are oxidized and converted into Zr0.80T0.20O2, B2O3 and SiO2, respectively, as shown in Fig. 6a,e. Zr0.80Ti0.20O2 partially melts and forms a relatively dense layer in the central ablated surface. However, the evaporants with low melting point, such as SiO2 and B2O3, escape from the oxide layer leaving evacuation channels: the holes, as shown in Fig. 6a. (2,500 °C), At a higher ablation temperature crystals connected by the melt become larger evidently shrink as shown in Figure 6b. Its cross  Zr0.8Ti0.2C0.74B0.26  the Zr0.80Ti0.20O2  and the holes section shows  04-014-0362: Carbon  (200)  (111)  (220)  (222) (311)  (002)  (100) (101)  (112)  (111)  (220)  (311)  04-001-2753: ZrC  0.6 4.75  0.7  0.8  0.9  1.0  4.75  4.70  4.65  4.60  4.55  - (111  )  - (311  )  (200)  [011]  4.50  a L  t t  i  c  e  c  n o  s  t  n a  t  (  Å  )  Atom ratio  of C: Zr of pure ZrC (%)  4.70  4.65  a L  t t  i  c  e  c  n o  s  t  n a  t  (  Å  )  4.60  4.55  4.50  b  a  c  Zr(Ti):C≈  1  Zr(Ti):C>1  Zr(Ti):C(B)≈  1  Atom ratio of metal: carbon (%)  Cross-section  Top surface  10  20  40  30  Zr0.8Ti0.2  Boron  Carbon  50  60  70  80  2 (♫)  00-029-1129: SiC  - (111  )  - (311  )  (200)  [011]  a  b  e  c  d  f  Figure 5 | Phases and structural changes in the Zr-Ti-C-B carbide during the PC process. (a) XRD results of top surface and cross-section of composites.  Black curve is from cross-section which is mainly composed of ZrC, carbon and a small amount of SiC. Red curve is from top surface which is mainly  composed of ZrC and SiC. However, TiC and diborides have not been observed due to the substitution of 20 at. % Zr by Ti and the vacancy filling of  boron atoms. (b) Variation of  lattice constant  in carbide crystal. The measurement of  lattice constant  is from the Rietveld refinement of XRD results  (Methods section). Red dots representing lattices of pure ZrC are from ref. 44. Columns are the lattice constants of carbide in this work. (c) Schematic  representation of distribution of boron atoms in Zr0.8Ti0.2C(1 \\x00 x) through defect channels. (d) Schematic representation of boron atoms filling in the carbon  vacancies. (e) High resolution transmission electron microscope (TEM) image of Zr0.8Ti0.2C(1 \\x00 x). The inset  in e is the electron diffraction pattern of  Zr0.8Ti0.2C(1 \\x00 x). (f) High resolution TEM image of Zr0.8Ti0.2C0.74B0.26 from the focused ion beam (FIB) samples. (Methods section and Supplementary  Fig. 3). The inset  in f  is the electron diffraction pattern of Zr0.8Ti0.2C0.74B0.26. Scale bar, 5 nm.  NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  ART IC L E  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications  5          \\x0c', 'A RT IC L E  NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  the porous morphology under the dense surface, as shown in Fig. 6f. Obviously, three different layers can be observed on the cross-section: porous external layer, intermediate layer and dense inner layer, possibly due to the evaporation of oxides and  thermal gradient perpendicular to oxide layer. For instance, higher temperature would result in more severe evaporation at the position closed to the external surface. However, the size of pores in the layers decreases significantly from the external layer  a  c  f  Surface  Resin  g  Surface  b  Zr Ti Si O B  e  )  .  u  .  a  (  y  t  i  s  n e  t  n  I  3,000°C  2,500°C  2,000°C  SiO2  ZT  Zr0.80Ti0.20O2  2,000°C  2,500°C  3,000°C  10  20  30  40  50  60  70  80  2 (♫)  Oxides  Carbide  d  )  %  . t  a  (  n o  i  t  c  a  r  F  80  70  60  50  20  15  10  5  0  External layer  Intermediate layer  Inner layer  Zr  Ti  O  B  Si  Oxides  Carbide  Figure 6 | Microstructure and phases of ablated surface and cross-section. (a) Central surface ablated at 2,000 °C. (b) Central surface ablated at 2,500 °C. (c) Central surface ablated at 3,000 °C. (d) Concentration of elements in central ablated surface. (e) XRD results of whole ablated surfaces. ZT is zirconium titanate having a-PbO2 structure45. (f) Cross-section morphology of central ablated point at 2,500 °C (back scattered electron images), and the  associated distribution of elements. The inset in (f) shows a higher magnification of 3,000 °C. Scale bar in a,b,c, 20 mm. Scale bar in f,g, 40 mm (inset, 2 mm).  inner layer. (g) Cross-section morphology of central ablated point at  6  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications        \\x0c', 'NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  ART IC L E  to the inner layer. Especially, the inner layer composed of the grain skeletons and amorphous phases displays a very dense morphology (see inset in Fig. 6f). It is inferred that the dense surface and inner layer act as a barrier to resist oxidation and result in reduced loss of oxide and the best ablation resistance shown in Fig. 1. At 3,000 °C, an almost fully dense oxide layer, mainly composed of Zr0.80Ti0.20O2, zirconium titanate and SiO2, is formed as shown in Fig. 6c,e. In addition, the XRD results suggest the presence of more amorphous phases, as the ablation 3,000 °C, temperature increases from 2,000 to due to the quenching of more liquid-solid phases of oxide layer at the end of ablation test (some oxides under the surface may remain solid at 3,000 °C within the limited ablation time, due to the thermal barrier provided by the Zr-O-Ti ceramic system33). Meanwhile, the sealing of the oxides can be attributed to these melts having a relative lower viscosity. However, such melts seal the defects (holes and cracks arising from the ablation) and protect the carbon matrix well, causing greater loss of the oxides by the scouring of hot high-speed gas as well, which is confirmed by the relatively low weight gains shown in Fig. 1 and the corrosion pores beneath the dense surface shown by the cross-section morphology of the sample after 3,000 °C ablation in Fig. 6g. However, a relatively dense inner layer located at the interface between oxide layer and carbide acted a barrier to the diffusion of oxygen, though some pores occurred in a thicker oxide layer formed with the temperature increasing from 2,500 to 3,000 °C. In addition, micro-cracks occurred on the cross-sections due to the thermal shock during the ablation test. The relatively integrated interfaces between oxide layers and carbides experiencing 2,500 and 3,000 °C ablation tests, showing a compact morphology without separation, indicate a good adhesion of oxide layers to substrates. Notably, it is believed that the dense Zr0.80T0.20O2 layer has effectively retained the boron and silicon and extended their consumption time, according to their residual contents shown in Fig. 6d,f.  Discussion  the Zr0.80Ti0.20O2  The carbide presented here exhibits superior ablation response compared to current UHTCs. This can be attributed to the following facts. First, a relatively dense oxide layer with a suitable viscosity plays a key role in resisting the ablation of extremely hot gas. Generally, loose scale with a very high viscosity (for example, no melt), or a liquid layer with a very low viscosity (for example, freely flowing), provides poor protection because of the tendency for detaching or splashing of oxides29, respectively. In this work, Zr0.80Ti0.20O2 displays a durable and robust structure comprising grain skeletons and liquid phases as shown in Fig. 6, which effectively decreases the loss of oxides, and the dense oxide layer displays good sealing protection even at 3,000 °C. In particular, the doping of Ti as the second main phase led to the formation of during the ablation test and decreased the viscosity of pure ZrO2 melt due to the lower melting point of TiO2 (1,843 °C) than that of ZrO2 (2,715 °C), which conferred a self-healing ability on the oxide layer, instead of a porous layer. Meanwhile this evidently decreased the vaporization and loss rate of oxides, compared with the conventional ZrB2-SiC or ZrC-SiC ceramic systems. For instance, the vaporization rate is 0.23 mm s \\x00 1 at 2,227 °C, whereas (VR) of TiO2 the VR of SiO2 (207 mm s \\x00 1) is around 900 times higher than that of TiO2 temperature34. at the same This suggests that oxidation of Zr0.8Ti0.2C0.74B0.26 will result in much less loss of oxide because of vaporization and good adhesion to the substrate shown in Fig. 6g even at 3,000 °C, compared to the ZrO2-SiO2 system from the oxidation of ZrB2-SiC and ZrC-SiC. Second, the good ablation response is due to the low oxygen permeability  the Zr0.80T0.20O2  (Zr0.80Ti0.20O2)  Zr0.8Ti0.2C0.74B0.26  (OP) of the oxide layer. In this work, the Zr0.80Ti0.20O2 layer on the quaternary carbide has effectively caught the B2O3 and SiO2 low OP (OPB2O3 ¼ 8.6 \\x02 10 \\x00 12 g cm \\x00 1 s \\x00 1 which has a very 1,000 °C, OPSiO2 ¼ 3.2 \\x02 10 \\x00 15 g cm \\x00 1 s \\x00 1 1,000 °C)35 at at and, to a certain extent, can prevent the fast diffusion of oxygen atoms into at different ablation temperatures. Moreover, it is believed that the intrinsic oxygen diffusion coefficient (ODC) of layer is lower than that (1.12 \\x02 10 \\x00 13 m2 s \\x00 1, of the pure ZrO2 due to the lower ODC of TiO2 1,800 °C) (ODCZrO2 ¼ 1. 16 \\x02 10 \\x00 12 m2 s \\x00 1, at compared with ZrO2 1,800 °C)36. at The lower oxygen permeability of Zr0.80Ti0.20O2 leads to less formation and loss of oxides, and further improves the ablation resistance of carbide, though the melting point of oxide is lower than 3,000 °C. Consequently, it is believed that our carbide displaying a better ablation resistance than the conventional ceramics can increase the survival time of the components in extremely oxidizing environments up to 3,000 °C. In addition, the lower boron content in Zr0.8Ti0.2C0.74B0.26 compared with that in ZrB2 leads to a lower mass loss and fewer defects such as pores and cracks originating from the evacuation channels of B2O3 at higher temperature, which is significantly beneficial to the ablation response. Third, it also benefits from the gradual transition from ceramic to carbon composite. A dense and high volume of ceramic is more conducive to forming a dense oxide layer and ensured the minimum damage to the underlying carbon matrix from the extreme hot and oxidizing gas. In conclusion, we have designed a carbide assembled by solid solution and atomic diffusion during PC and RMI. Importantly, the ceramic layer displays better ablation resistance (eg, a rate of material loss over 12 times better than conventional ZrC at 2,500 °C) under oxidizing environments up to 3,000 °C relative to existing common UHTCs and high temperature composites. More broadly, this work provides a platform for building a series of such UHTCs (eg, A (M) CxBy), where A and M are the main atoms (transition metals, IV) and the substitution atoms (transition metals, IV/V), respectively. For instance, Hfx (Zry1/Tiy2/Tay3)C0.8B0.2, Zrx (Tiy1/Tay2)C0.8B0.2 and TiC0.8B0.2 and so on, can be built, according to similar methods. To increase ablation resistance across a wide range of temperatures, a carbide’s oxides layer that owns low VR, OP, good self-healing ability and adhesion strength to the substrate can be achieved through the doping contents of different substitution atoms (yi) of ceramic with a variety of melting points. Consequently, we develop an effective means of fabricating ablation resistant UHTCs. Moreover, these ceramics can be fabricated into powders, bulk materials and layers to extend their application. For instance, in addition to the potential use in hypersonic vehicles, it is expected that they can be used as the nozzle throats and diffusers for reusable rockets, which requires a very low ablation rate to extend the lifetime for recyclability at low cost. Other potential uses may include the hot section components in re-entry spacecraft, defence army, gas turbines and nuclear areas and so on.  Methods  Materials and preparation. Supplementary Fig. 1 shows a schematic of the sample preparation process. T700 Polyacrylonitrile-based carbon fibres were employed as the reinforcement and fabricated into a needled integrated preform (NIP). The NIP was fabricated by a three dimensional needling technique, built up by repeatedly overlapping layers of 0° non-woven C-fibre cloth (A in Supplementary Fig. 1a), a chopped fibre web (B in Supplementary Fig. 1a), and 90° non-woven fibre cloth with needle-punching step by step. The bulk preforms were densified to a porous C/C composite of 1.0-1.3 g cm \\x00 3 density using pyrocarbon deposited by CVI using CH4 and H2 gases at 900-1,000 °C. The open porosity of the composites with NIP ranged from 39.8 % to 28.8 %. For instance, a sample having a density of 1.16 g cm \\x00 3 and an open porosity of 34.3 %  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications  7  \\x0c', 'A RT IC L E  NATURE COMMUNICATIONS | DOI: 10.1038/ncomms15836  possessed a modal pore diameter of 42.0 mm shown in Supplementary Fig. 6. Because non-stoichiometric carbide tends to form when carbon reacts with Zr-Ti melt (Supplementary Fig. 2), a Zr0.8Ti0.2C(1 \\x00 x) carbide was introduced into the in argon at 1,800-2,000 °C for 0.5-2 h. C/C composite by RMI In this step, an optimized ratio of powders mixed as 80 at. % Zr-20 at. % Ti was used29. The the temperature of 1,800-2,200 °C ranged viscosity of 80 at. % Zr-20 at. % Ti at from 4.55 \\x02 10 \\x00 3 to 3.06 \\x02 10 \\x00 3 Pa s as shown in Supplementary Fig. 7. The molten metal was drawn along the carbon fibres by capillary forces, reacting with the previously deposited pyrocarbon to form a carbon/carbide composite. The infiltrated depth and the volume of carbide are dependent on the mass of the melt and the volume and size distribution of the open pores in the C/C composite. Finally, a carbide layer was formed on carbide-based part by PC process (Supplementary Fig. 1c). During the PC process, B4C, B2O3, SiC, Si, carbon and some catalysts such as Al2O3 were the raw powders. The carbide-based part was packed by powders and placed in a graphite crucible. At 1,600-1,800 °C, these powders react with Zr0.8Ti0.2C(1 \\x00 x) and carbon fibres and pyrocarbon to form the Zr0.8Ti0.2C0.74B0.26 ceramic and SiC, respectively. In addition, to achieve relatively low porosity, the carbon-based part of sample was further densified using CVI. For comparison of ablation performance, C/C-Zr0.8Ti0.2C, ZrTiC-SiC and C/C-SiC composites were fabricated by RMI. In addition NIP processed C/C composite having a density of 1.7 g cm \\x00 3 was densified by CVI. ZrB2-SiC ceramics were prepared by SPS. More details about the fabrication were referenced Supplementary Note 1.  Ablation testing. The ablation behaviours of the samples were evaluated using an oxyacetylene flame (Supplementary Movie 2). During the test, the specimen, having a size of | 30 \\x02 15 mm, was exposed to the flame. The flow rates and pressure of oxygen were respectively 1.96 l s \\x00 1 and 0.400 MPa, and those of acetylene were 0.696 l s-1 and 0.095 MPa, respectively. The normal combustion ratio of oxygen and acetylene is 1.5, according to 2C2H2 þ 3O2 ¼ 4CO þ 2H2O, and in this work the extra oxygen ensured a sufficient combustion of acetylene and established an extreme oxidizing scenario. During the test, an optical pyrometer indicated that the highest temperature of the central ablated surface reached about 3,000, 2,500 and 2,000 °C at the distances of 10, 20, and 30 mm between the torch nozzle and sample surfaces, respectively. The heat fluxes measured by a watercooled heat flux sensor at 3,000, 2,500 and 2,000 °C are 5.62, 3.86, 2.57 MW m \\x00 2, respectively. The inner diameter of the oxyacetylene gun tip was 2 mm. The linear ablation rate (LAR) and mass ablation rate (MAR) were calculated according to _l ¼ Dl= Dt \\x01 S Þ and _m ¼ Dm= Dt \\x01 S Þ respectively, where _l refers to LAR and _m to MAR; Dl and _m are the decrease in length and the mass loss of specimen, respectively, and Dt is the ablating time. S is the ablation area. Ablation rate is averaged over three specimens. The tests lasted 60 and 120 s.  ð  ð  Characterization methods. 3D X-ray computed tomography (CT) was conducted using a Zeiss Xradia Versa 520 X-ray microscope at the Henry Moseley X-ray Imaging Facility (HMXIF, Manchester, UK). The accelerating voltage and current of the X-ray tube were set as 140 kV and 72 mA, respectively. Each scan comprised 1,601 radiographs taken incrementally over a rotation angle of 360˚ . A 3D volume rendering of the sample was created from the virtual slices in AVIZO software. The first 100 slices and last 100 slices were cropped due to low image quality presumably as a result of the cone-beam geometry. Pores, carbon and ceramics were segmented out using the top hat method37 (Supplementary Movie 1). XRD experiments were carried out on a Panalytical MPD system using Cu radiation. The voltage and current were 45 kV and 40 mA, respectively. For the analysis of phases, the 2y scan range was 10-90°, scanning resolution was 0.05° per step. For the measurement of crystals structure, the samples were scanned between 10-90° 2y, at a step-width of 0.02° and scan speed of 0.5° \\x01 per min. Rietveld refinement was carried out with the program Maud using a pseudo-Voigt profile function38,39, and Rw was less than 9.6%. A super probe electron probe microanalysis system (EPMA, JEOL, Jxa8230) was used to detect the content and distribution of elements. High resolution transmission electron microscope (HRTEM) images and selected area electron diffraction (SAED) patterns were obtained with an FEI Talos F200A microscope equipped with an X-FEG electron source. TEM samples were prepared by focused ion beam (FIB, FEI Quanta 3D) using the in-situ lift-out technique on cross-sections of the samples40. The morphology of samples was studied by scanning electron microscopy (SEM, FEI., NOVA Nano230). Bulk density was measured according to the Archimedes method41. The 3D surface profile was conducted using the 3D Optical Microscopy (Bruker Contour Elite 3D Optical Microscope). The open porosity was measured using the boiling water method according to the ASTM Standard C20-00, and the pore size distribution of NIP samples was investigated using mercury porosimetry (Quantachrome, Pore Master 60), according to ISO 15901-1.  Data availability. 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Preparation of SiC/SiC composites by hot pressing, using Tyranno-SA fiber as reinforcement. J. Am. Ceram. Soc. 86, 26-32 (2003). 42. Chen, Z. K., Xiong, X., Li, G. D. & Wang, Y. L. Ablation behaviors of carbon/carbon composites with C-SiC-TaC multi-interlayers. Appl. Surf. Sci. 255, 9217-9223 (2009). 43. Hao, Z. H., Sun, W., Xiong, X., Chen, Z. K. & Wang, Y. L. Effects of Ti/Al addition on the microstructures and ablation properties of Cf/C-MoSi2-SiC composites. J. Eur. Ceram. Soc. 36, 457-464 (2016). 44. Pierre, V. Material Phases Data System (MPDS) (Springer, 2014). 45. Troitzsch, U. TiO2-doped zirconia: crystal structure, monoclinic-tetragonal phase, transition, and the new tetragonal compound Zr3TiO8. J. Am. Ceram. Soc. 89, 3201-3210 (2006).  Acknowledgements  We thank National Basic Research Program of China (No.2011CB605805), China  Postdoctoral Exchange Fellowship Program (20140012) and National Natural Science  Foundation of China (51602349) for funding of this work. The authors would also like to  thank the Henry Moseley X-ray Imaging Facility (HMXIF, Manchester, UK) for x-ray  tomography measurement  funded by the EPSRC through grants EP/M010619,  EP/K004530, EP/F007906, EP/F028431. We are indebted to Han Liu in University of  Manchester for the assistance with calculation of  the thermodynamics.  Author contributions  X.X., P.X. and Y.Z. proposed and designed the project. Y.Z. and D.W. developed the  optimal composition of ceramic and fabricated the ceramic composites and the analyses  including SEM, TEM, XRD and EPMA. M.B. carried out the FIB and M.S. conducted the  TEM operation and analyses. W.S. and D.W. conducted the ablation test and  measurement of property. X.Z. and P.W. carried out  the 3D X-ray tomography  and the analyses. Y.Z. wrote the paper with input  from all authors, and P.W., X.X. and  P.X. refined the paper. All authors contributed to the interpretation of  the results.  Additional  information  Supplementary Information accompanies this paper at http://www.nature.com/  naturecommunications  Competing interests: The authors declare no competing financial  interests.  Reprints and permission information is available online at http://npg.nature.com/  reprintsandpermissions/  How to cite this article: Zeng, Y. et al. Ablation-resistant carbide Zr0.8Ti0.2C0.74B0.26 for oxidizing environments up to 3,000 °C. Nat. Commun. 8, 15836  doi: 10.1038/ncomms15836 (2017).  Publisher’s note: Springer Nature remains neutral with regard to jurisdictional claims in  published maps and institutional affiliations.  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To view a copy of  this  license, visit http://creativecommons.org/  licenses/by/4.0/  r The Author(s) 2017  NATURE COMMUNICATIONS | 8:15836 | DOI: 10.1038/ncomms15836 | www.nature.com/naturecommunications  9  \\x0c']"
},{
  "_id": 16,
  "PDF": "Anti-oxidation modification behaviors and mechanisms of ZrB2 phase on Si-based ceramic coatings in aerobic environment with wider temperature region.pdf",
  "Text": "['Journal of Alloys and Compounds 769 (2018) 387e396  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : h t t p : / / w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  Anti-oxidation modification behaviors and mechanisms of ZrB2 phase on Si-based ceramic coatings in aerobic environment with wider temperature region  Xuanru Ren*, Xiaojun Sun, Wenhao Wang, Hongsheng Mo, Peizhong Feng**, Li Tong Guo***, Ziyu Li  School of Materials Science and Engineering, China University of Mining and Technology, Xuzhou, 221116, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 30 April 2018 Received in revised form 22 July 2018 Accepted 25 July 2018 Available online 27 July 2018  Keywords:  ZrB2 Powders Coating Liquid phase sintering Anti-oxidation modification Zr-oxidesskeleton  To improve the modification effect of ZrB2 phase on Si-based ceramic coatings in aerobic environment with wider range of temperature, ZrB2 was utilized to modify the SiC coating by the technique of liquid phase sintering. The modification behaviors of ZrB2 phase on oxidation inhibition ability of SiC coating were investigated in wider range of temperature through the TG oxidation tests (room temperature1773 K). When the synthetic temperature is 1700 \\x0e C, pure phase ZrB2 powders were synthesized through the way of carbothermal reduction, the average particle size of which is 531 nm. The SiC coating modified by ZrB2 is comprised of ZrB2 and SiC phases, the thickness of which is about 200 mm. With the modification of ZrB2 phase, the oxidation resistance of the SiC coating is significantly enhanced, especial the temperature region below 1200 \\x0e C, initial mass loss temperature (about 260 \\x0e C) leading to the delay of and slowing down mass loss rate (reduced by about 67%) in the fastest mass loss area. While the final mass loss of the coated graphite substrate decreased from 18% to 5%. Owing to the formation of the protective B2O3, the fast weight gain temperature region (from 700 \\x0e C to 1240 \\x0e C) in the TG curve of the ZrB2 powders effectively compensate the weakness of the SiC coating in this temperature region. With the movement of the increasing SiO2 glass layer to cover the surface of the coating, Zr-oxides consisted of ZrO2 and ZrSiO4 were gradually peeled into tiny oxide particles to form Zr-oxides-rings. The refractory Zroxides-rings and the unpeeled large Zr-oxides-obstacles construct skeleton inlayed in the SiO2 glass layer, demonstrating the ability of restricting the growth of microcrack, which is responsible for the remarkably enhanced anti-oxidation modification effect of ZrB2 phase. © 2018 Elsevier B.V. All rights reserved.  1.  Introduction  As a result of the superior strength retention, low density and coefficient of thermal expansion [1e3], carbon structural materials are regarded as promising candidate materials in aerobic environments with super high temperature as applications of aircraft and aerospace. However, C is apt to be severely oxidized in an aerobic environment when the temperature is above than 500 \\x0eC, whose mass loss will damage their structure, thus greatly restricting their  * Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: XuanruRen@163.com (X. Ren), litongguo@cumt.edu.cn (L.T. Guo).  https://doi.org/10.1016/j.jallcom.2018.07.300 0925-8388/© 2018 Elsevier B.V. All rights reserved.  fengroad@163.com (P. Feng),  broad applications [4e6]. To solve the obvious defects of carbon structural materials, the anti-oxidation coating technology [7e9] are deemed to be the most effective way to restrict the erosion of oxygen to the carbon matrix. Among the numerous coating materials, ceramics chemicals containing Si elements are recognized as the most valid antioxidation coating components [10e13], such as SiC [14], MoSi2 [15], CrSi2 [16] and TaSi2 [17], which are able to provide valid role of inhibition of oxygen corrosion. The key of the oxidation protective behaviors of the ceramics chemicals containing Si elements is the formation of silicate glass layer. The silicate glass layer is a kind of oxidative protective layer with the ability of self-healing above 1200 \\x0eC, which is capable of reducing the erosion of oxygen to the carbon matrix through the way of sealing the defects and decreasing the permeability of oxygen. However, owing to the excellent thermal structural properties, the demand of the carbon  \\x0c', \"388  X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  structural materials is extensive, which results in the need in very volatile aerobic environments. Thus, the carbon structural materials not only need to be applied in a constant temperature environment, but also in an environment with variable temperature changes. However, besides the finite oxidation protective temperature region, the application of Si-based ceramic coating is restricted because in which cracks and bubbles will be generated after serving a long time at an environment with variable temperature changes, which actually exist as“diffusion paths”of oxygen to C/C composites, resulting in the failure of the C/C composites [18,19]. It is extremely urgent to enhance the oxidation resistances of the Si-based ceramic components under varied aerobic environments. Up till now, combining the outstanding properties of the modifiers, the Si-based coating modification technology has attracted much attention, which is able to enhance the oxygen inhibition ability of the Si-based ceramic coatings on its inherent characteristics. Among the applied modifiers, thanks to the high melting point, promising ability of oxidation resistance and excellent thermal shock resistance [20e25], the ZrB2 ceramic, one of the ultra-high temperature ceramics (UHTC) borides [26e32], presents great potential modification effect on the Si-based ceramic coatings. When the ZrB2-SiC ceramic was exposed in an aerobic environment, multi-component oxides consisted of ZrO2, SiO2, ZrSiO4 and B2O3 will be generated. The ZrO2 and ZrSiO4 are well-known as stable super high temperature resistant materials, whose melting points are 2700 \\x0eC and 2500 \\x0eC, respectively. While the B2O3 is not only a suitable medium temperature protection material below 1200 \\x0eC, but also an important source of material to react with SiO2 to form borosilicate glass layer with better stability and oxidation resistance. Hence, the ZrB2 presents the great potential to modify the anti-oxidation ability of the Si-based ceramic components in an aerobic environment with wider range of temperature. In our previous work, utilizing graphite, B2O3, ZrO2 and Si as raw materials, we have synthesized ZrB2-SiC/SiC coating on carbon/ carbon composites through in-situ reaction method [33], and under the protection of the SiC coating modified by the ZrB2, the mass loss of the C/C composites is only 0.22% at 1773 K for 550 h, exhibiting excellent modification effect. However, although microstructure of ZrB2 phase and its contents in coating are essential for its antioxidation modification effect on the Si-based ceramic, the both factors are actually incontrollable in the preparation process of the in-situ reaction method, which limits the scope of the modification effect of ZrB2 phase in aerobic environment with wider range of temperature. Recently, to overcome the shortcomings of the in-situ reaction method, in our former work, a liquid phase sintering technique was designed to prepare the SiC coating modified by TaB2 phase [34], actualizing the control of the microstructure of TaB2 phase as well as its contents in coating, which is suitable to be used to modify the Si-based coatings with other anti-oxidation components. Hence, in this work, to improve the modification effect of ZrB2 phase on Sibased ceramic in aerobic environment with wider range of temperature, the technique of liquid phase sintering was utilized to modify the SiC coating by ZrB2 phase. Furthermore, the oxidation protective behaviors and mechanisms of the SiC coating modified with ZrB2 phase at dynamic aerobic environment with wider range of temperature from room temperature to 1773 K were investigated.  1.1.  Experimental procedures  SiC coating modified with ZrB2 phase was synthesized by liquid phase sintering technique, which consisted of powder synthesis of ZrB2 and coating preparation.  Fig. 1 (a) shows the synthesis of the ZrB2 powders. The original materials were composed of Si (9e20 wt%, Jiuling Smelting Co., Ltd., Shanghai, China), graphite (4e18 wt%, Carbon Plant., Xi'an, China), B2O3 (10-30 wt%, Tianli Chemical Reagent Co., Ltd., Tianjin, China) and ZrO2 (40-60 wt%, Guoyao Chemical Reagent Co., Ltd., Shanghai, China) powders. Firstly, the raw powders were thoroughly mixed in a clean ball mill for 2 h. Then, the homogeneous raw powders mixture was heat-treated with the protection of argon atmosphere, the temperature of which were set as 1773 K and 1973K. The rate of the heat-treatment was 5e10 K/min, and the holding time at the corresponding temperature was 2 h. During the process of heat treatment, as shown in Eq. (1), carbothermal reduction reaction would occur, and finally the ZrB2 phase would form by reducing ZrO2 using B2O3 and graphite.  ZrO2(s) þ B2O3 (s) þ5 C(s) / ZrB2 (s) þ5 CO (g)  (1)  (the Vsilica  sol: M other  raw materials ¼ 0.6e1.6 ml/g)  The inner SiC transition buffer coating was prepared on graphite (3 mm \\x02 3 mm \\x02 3 mm), whose process was described matrices previously [33]. The outer ZrB2-SiC coating overlaid on the surface of SiC coating was prepared by the method of liquid phase sintering, whose illustration was presented in Fig. 1(b). First, the synthesized pure ZrB2 powders (40 wt.%), graphite (7e21 wt.%) (Carbon Plant., Xian, China), SiC powders (15e30 wt.%) (China New Metal Materials Tec Co., Beijing, China), Si powders (10e28 wt.%) (Jiuling Smelting Co., Ltd, Shanghai, China), and silica sol (SiO2$ nH2O) (City Fire Crystal Glass Co., Ltd, Dezhou, China) were weighed as original materials for the synthesis of the original slurry. Second, the weighed raw materials were agitated one hour using a magnetic agitator to prepare mixed slurry. Third, the mixed slurry was brushed onto the surface of the inner SiC coating in order to form the pre-fabricated layer. The process of brushing repeat three times, and the samples were dried at 373 K for half an hour after each brushing. Finally, the samples placed in normal argon atmosphere were heat-treated to synthesize the outer coating. The heattreatment temperature was set as 2373 K with a holding time of 2 h and heating rate of 5e10 K/min. During heat treatment, since there would be the in-situ reaction among Si, silica sol and C powders, SiC would be formed as shown in reaction (2) and (3). In addition, as we know, the melting temperature of Si and silica sol are 1687 K and 1923 K, respectively, meaning that they would become the liquid state at the temperature of above about 2000 K during the heat-treatment. It can offer the required fluid factors for liquid phase sintering. Hence, the presence of Si powders and silica sol could largely improve the densification process of the coating.  Si (s) þ C(s) / SiC (s)  SiO2 (s) þ 3 C(s) / SiC (s) þ 2 CO (g)  (2)  (3)  The phase compositions analysis of the ZrB2 powders and the SiC coating modified with ZrB2 was examined using an X-ray diffractometer (XRD, Bruker D8 ADVANCE, BRUKER AXS, Germany). The surface element analysis of ZrB2 powders was characterized by X-ray photoelectron spectroscopy (XPS, Axis Ultra, Kratos Ltd., England). The Laser Particle Size Analyzer (Zen 3600, Malvern, England) was utilized to measure the particle size distribution of the synthesized ZrB2 powders. The Field Emission Scanning Electron Microscopy (FE-SEM, JSM-6700 F, JEOL, Japan) and energy dispersive spectroscopy (EDS) were performed together to analyze the morphology and elemental component of the powders and coating under various situations. High-resolution transmission electron microscopy (HR-TEM, JEM-3010, JEOL, Japan) was applied to measure the crystal structure features of the ZrB2 powders. The  \\x0c\", 'X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  389  Fig. 1.  Illustration of the synthesis of ZrB2 powders (a) and SiC coating modified with ZrB2 (b).  oxidation behaviors of the coating in aerobic environment with dynamic temperature regions from room temperature to 1773 K were carried out using a TG-DTA device (TGA, STA 449 F3, Netzsch, Germany). After the TGA tests, the samples were placed in electrical furnace for 50 h at 1773 K with air atmosphere to further analyze the oxidation characteristics of the coating.  2. Results and discussion  Fig. 2 shows TG-DSC curves of the ZrB2 precursor powders. The TG curve can be divided into four stages. The first stage is from room temperature to 110 \\x0eC, where the mass loss is negligible. The second stage is from 110 to 400 \\x0eC, over which the mass loss is greatest. In this stage, sharp endothermic peaks can be observed near 200 \\x0eC, causing by the dispelling of the absorbed water in ZrB2 precursor powders. The third stage is from 400 to 1400 \\x0e C, where the mass loss tendency is gentle. The fourth stage is from 1400 to 1500 \\x0eC, over which there is an mass loss causing by the carbothermal reduction reaction. At the same time, an exothermic peak exists near 1400 \\x0eC, indicating the beginning of the carbothermal reduction. In our previous work [33,34], 1500 \\x0e C and 1700 \\x0eC are two common temperatures to prepare boride powders and with the growth of temperature, precursor powders would react more adequately. Thus, in this experiment, we chose 1500 \\x0e C and 1700 \\x0e C as the reaction temperature to find a more suitable temperature for carbothermal reduction. The original powders are heat-treated at 1500 \\x0eC and 1700 \\x0eC respectively. The XRD patterns of the synthesized ZrB2 powders are shown in Fig. 3. At the temperature of 1500 \\x0eC, not only ZrB2 phase but ZrC phases can be seen in the XRD pattern, indicating ZrO2 and C had reacted during the carbothermal reduction reaction. Nevertheless, when the temperature rises to 1700 \\x0eC, except for a small amount of carbon, there is only ZrB2 phase can be seen in the XRD pattern, demonstrating that carbothermal reduction reaction of ZrB2 occurred adequately, which in1700 \\x0e C dicates the temperature of is more suitable for the  preparation of ZrB2 by the method of the carbothermal reduction reaction. Fig. 4(a) describes the SEM micrographs of the ZrB2 powders, in which there are lots of fine clusters can be seen. By roughly calculating, the particle size of the majority of the synthetic ZrB2 powders is below 800 nm. Fig. 4 (b), the specific volumecumulative distribution curves of the ZrB2 powders, furthermore investigates the particle size of the synthetic ZrB2 powders. As can be seen, apart from the existence of a spot of larger particles in the powder, most of the powders distribute in a size range of 120e800 nm, which nearly accounts for 90% of the total ZrB2 powders. In addition, the average particle size of the synthetic powders is 531 nm, which is generally consistent with the result of SEM shown in Fig. 4 (a). The binding energies of the Zr4p, Zr3d, B1s, Zr3p3/2 and Zr3p1/2 peaks of synthetic ZrB2 powders are 28.1eV,  Fig. 2. TG-DSC curves of the ZrB2 precursor powders.  \\x0c', '390  X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  Fig. 3. XRD patterns temperatures.  of  the  synthesized  ZrB2  powders  heat-treated  at  different  180.5eV, 189.8eV, 330eV and 344eV, respectively. These binding energies are consistent with reported values by Aizawa et al. [36] for single-crystal bulk ZrB2, which are 27.5 eV (Zr4p2/3), 29.2 eV (Zr4p1/2), 178.9 eV (Zr3d5/2), 181.3 eV (Zr3d3/2), 187.9 eV (B1s), 329.9 eV (Zr3p3/2), 343.5 eV (Zr3p1/2), further illustrating the formation of ZrB2 phase. These results are basically consistent with the XRD analysis. For the sake of further exploration of the microstructure and crystal structure of the synthetic ZrB2 powders, TEM analyses were carried out. The Fig. 5 (a) shows the low magnification TEM micrograph of the ZrB2 powders. As can be seen, despite the fact that ZrB2 powders generally possess tiny sizes, the particle sizes of the ZrB2 powders are irregular and changes in the range of about 30 nme500 nm due to the agglomeration of the powders. There exist only a few larger size particles (700 nme1000 nm) in ZrB2 powders. Fig. 5(b) shows the high resolution TEM micrograph, further detecting the lattice fringes of the ZrB2 particles. In this micrograph, there are two kinds of atomic lattice spacing can be observed which are 2.86 Å and 3.57 Å by measurement, matching the (001) and (100) planes of ZrB2, respectively, which is consistent with the XRD results. The pattern of the ZrB2-SiC coating synthesized through liquid phase sintering method is shown in Fig. 6. It can be observed that the SiC coating modified with ZrB2 is only composed of ZrB2 and SiC phases, which indicates that during the process of liquid phase sintering, both of the ideal products are simultaneously formed. In addition, neither excess Si powders nor silica sol was detected after the process of heat-treatment, which indicates Si powders and silica sol adequately reacted, implying the liquid phase sintering method to prepare multiphase ceramic coatings is feasible. Fig. 7 (a) shows the cross sectional backscatter micrograph of the SiC coating modified with ZrB2. It can be seen that the thickness of the coating is approximate 200 mm. No obvious big hole or penetrating crack can be observed in the coating due to the effective combination between the double layers. Moreover, the visible microcrack is not existed in the cross-section of the coating either. It means that there is no path provided by microcracks to allow oxygen to diffuse, which will greatly improve the oxidation protection ability of the coating. For the above reasons, we can conclude that  Fig. 4. SEM micrographs synthetic ZrB2 powders.  (a), particles  size distribution(b), XPS spectrum (c) of  the  there is a huge prospect of using the method of liquid phase sintering to prepare the modified coating. Fig. 7 (b) shows the surface backscatter micrograph of the SiC coating modified with ZrB2, from which there are two kinds of particles can be observed. By EDS analyses (Fig. 7(c)), the two kinds of particles are SiC and ZrB2 phase, respectively, which are in accord with the XRD pattern in Fig. 6. The TGA curves of the coating which further illuminates the oxidation protection ability of the coating in air from room temperature to 1500 \\x0e C are shown in Fig. 8. As can be seen, the graphite the temperature of about 540 \\x0eC substrate starts to lose mass at under the protection of pure SiC coating, and the final mass loss of the coated graphite substrate reaches to 18%. While with the  \\x0c', \"X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  391  Fig. 6. XRD pattern of method.  the ZrB2-SiC coating synthesized by liquid phase sintering  shown in Fig. 10, when the temperature is over 700 \\x0eC, As the weight of ZrB2 rises rapidly which indicates the start of the oxidation reaction between ZrB2 and O2. Afterwards, the weight of ZrB2 increases with very rapid rate due to the formation of ZrO2 and B2O3. In addition, a weight gain peak can be seen when the temperature is about 1240 \\x0eC. The amount of ZrO2 and low melting B2O3 phase increases along with the rise of temperature. Owing to the inhibitory effect of B2O3 on oxidation, the increased amount of the range of 700 \\x0e Ce1240 \\x0eC explains the delayed initial B2O3 at mass loss temperature and reduced average mass loss rate in the fastest mass loss area of the pure SiC coating modified by ZrB2 phase. When the temperature is higher than 1300 \\x0eC, the B2O3 who has a low boiling point starts to evaporate, which leads to mass loss of ZrB2 powders. However, the phenomenon of mass loss can't be seen in Fig. 9 in ZrB2-SiC/SiC coating after 1300 \\x0eC, which is brought about by the inhibitory effect of SiO2 glass on the volatilization of B2O3 through the formation of borosilicate glass layer. Fig. 11 shows XRD pattern of the ZrB2-SiC coating after oxidation. As can be seen, after oxidation at 1773 K, there are some oxides formed on the surface of the coating, including ZrSiO4, SiO2 and ZrO2. Therefore, it can be concluded that a Zr-Si-O glass layer is finally formed on the surface of the coating. The ZrO2 and ZrSiO4 are well-known as stable super high temperature resistant materials, whose melting points are 2700 \\x0eC and 2500 \\x0eC, respectively. With the multiple oxides in the compound glass layer, the oxidation protective ability of the coating is capable of improving further [35e37]. Fig. 12 displays the surface backscatter micrograph of SiC coating modified with ZrB2 after TGA test, in which we can see the surface microstructures of the coating after TGA oxidation. As shown in Fig. 12(a), many white phase particles scatter on the surface of the coating after TGA test, which indicates the formation of Zr-oxides by the EDS spot analyses. Moreover, the grain boundary of the black SiC particles are clear and complete, indicating the continuous SiO2 glass layer has not formed due to the limited oxidation time of TGA test. All the same, the phenomenon of “halation” in SiO2 glass layer can be observed. To clearly observe the microstructure of the “halation”, the Fig. 12(a) is magnified as shown in Fig. 12 (b). As can we see, though the SiO2 glass layer has formed, but not enough to cover the gaps among the SiC grains.  Fig. 5. TEM micrograph (a) and high-resolution TEM micrograph (b) of the synthetic ZrB2 powders.  modification of ZrB2, the start mass loss temperature of the graphite substrate rises to about 800 \\x0e C and the final mass loss of the graphite substrate drops to 5%. The start mass loss temperature of the graphite substrate rises by about 350 \\x0e C and the final mass loss of the graphite substrate reduces by 13% with the modification of ZrB2, which indicates that the ZrB2 phase has promising oxidation protective modification ability. The mass loss rate curves of the coating during TGA test from room temperature to 1500 \\x0eC are shown in Fig. 9, from which we can get three important points. The first point that needs to be explained is the initial mass loss temperature of the graphite substrate rises by about 260 \\x0e C with the modification of ZrB2. The loss area with the addition of ZrB2 is \\x002 \\x02 10 second point is that the average mass loss rate in the fastest mass \\x003 mg cm \\x002 s \\x001, which reduced by about 67% compared with the \\x006 \\x02 10 \\x003 mg cm \\x002 s \\x001 of the pure SiC coating, causing by the formation of B2O3. The melting temperature of B2O3 is about 450 \\x0eC, which means it will become liquid with the increase of temperature and form a protection layer to protect C substrate. The final point is that the range of the fastest mass-loss zone under the protection of the 40%ZrB2-SiC coating is much shorter than the pure SiC coating. Thus, it can be seen that the oxidation protective ability of the coating in the area of middle temperature (600\\x0eC-1200 \\x0e C) is significantly increased with the modification of ZrB2, which effectively compensates the weakness of the anti-oxidation ability of SiC coating. In order to further reveal the modification mechanisms of ZrB2 phase on Si-based ceramic coating, TG test of the synthetic ZrB2 powders was conducted in air from room temperature to 1500 \\x0eC.  \\x0c\", '392  X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  Fig. 7. Cross section (a) and surface (b) backscatter SEM micrographs of the SiC coating modified with ZrB2; (c) Spot EDS analyses.  However, owing to the fluxility of SiO2 glass layer at superhigh oxidation temperature, the white Zr-oxides are peeled away by the fluid SiO2 glass and become smaller particles which can inlay in the SiO2 glass layer. Fig. 13 (a) shows backscattered SEM micrograph of the ZrB2-SiC coating after oxidation at 1773 K for 50 h. It can be seen that after oxidation at 1773 K for 50 h, a dense compound glass layer has formed on the surface of the ZrB2-SiC coating, in which many white particles inlay. Moreover, there is no clear long crack exists. Fig. 13(b) shows the magnification of part A in Fig. 13(a). It can be seen that besides the inlayed large white oxides, an interesting  phenomenon of Zr-oxides-rings can be seen in the glass layer. By EDS analyses (shown in Fig. 13(c)), it can be concluded that the white particles are Zr-Si-oxides consisted of Zr, Si and O elements, indicating the generated ZrO2 and ZrSiO4, while the grey glass layer is SiO2 glass layer consisted of Si and O elements. The Zr-oxidesrings are formed by the movement of the peeled Zr-oxides with the fluid SiO2 glass layer, which will expand the diffusion area of the Zr-oxides. Owing to the super high melting temperature of the Zroxides, the Zr-oxides rings are actually kinds of super-high temperature skeletons inlayed in the SiO2 glass layer, while the inlayed  Fig. 8. TGA curves of the coatings in air from room temperature to 1500 \\x0e C.  Fig. 9. Mass loss rate curves of the coating during TGA test from room temperature to 1500 \\x0e C.  \\x0c', 'X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  393  Fig. 10. TG curve of the ZrB2 powders in air from room temperature to 1500 \\x0e C.  unpeeled large oxides are super-high temperature obstacles, which will increase the stability of the compound Zr-Si-O glass layer. With the existence of the Zr-oxides skeletons and large obstacles, the phenomena of crack turning, thinning and pinning can be seen when the cracks pass through the glass layer, which indicates the Zr-Si-O glass layer possesses the ability of consuming the energy of the cracks and hindering the elongation of cracks. The cross-sectional backscatter micrograph of the multilayer coating after oxidation is depicted in Fig. 14(a) and (b). Firstly, there forms a thin Zr-Si-O glass layer on the surface of the coating after oxidation, and no obvious hole can be seen inside the coating, which indicates the effective shielding of the glass layer on oxygen. Secondly, it can be seen although a few microcracks unavoidably exists in the surface of Zr-Si-O glass layer, there is no penetrating crack inside coating or graphite substrate, which illustrates that the Zr-Si-O glass layer can hinder the extension of microcracks. Thirdly, the associativity among Zr-Si-O glass layer, coating and graphite substrate is compact after oxidation and no spalling is observed. Fig. 15 illuminates the sketch map of oxidation protection  Fig. 11. XRD pattern of the SiC coating modified with ZrB2 after oxidation.  Fig. 12. (a) Surface backscatter micrograph of the SiC coating modified with ZrB2 after TGA test; (b) magnification of part A in (a).  mechanisms of the SiC coating modified with ZrB2. As shown in Fig. 15(a), before oxidation, there was a double-layer coating on the C/C composites, which is the first barrier to isolate the carbon matrix from the oxygen. The quality of the coating determines the oxidation protection performances of the coating at this stage. As shown in Fig. 15(b), at the very first of oxidation process, the outer ZrB2-SiC coating exposed in the air atmosphere firstly and reacted with oxygen. In this stage, some oxides formed including SiO2, ZrO2 and B2O3, whose reaction equations are shown in Eq. (4) Eq. (6). In the temperature below 1200 \\x0e C, with the prolongation of oxidation time and increase of temperature, besides the anti-oxidation protective ability provided by the coating itself, the formed B2O3 will become liquid layer, which is capable of resisting the erosion of oxygen to the carbon matrix. While when the temperature is above 1200 \\x0eC, the formed SiO2 and ZrO2 would provide corresponding role of anti-oxidation [38]. At this stage, with the increase of temperature, more SiO2 will be formed, which gradually became liquid state and began flowing to cover the surface of the coating. By the way of flow, the generated SiO2 glass will seal the defects and reducing the oxygen permeability, thus constructing another protective system [39]. Moreover, the generated SiO2 glass is able to inhibit the volatilization of B2O3 by the way of forming borosilicate glass layer. Furthermore, the reaction occurred between the formed ZrO2 and SiO2, generating ZrSiO4 phase as shown in Eq. (7), which further enhance the oxidation resistance of the designed coating [40]. As shown in Fig. 15(c), with the movement of the increasing SiO2 glass layer to cover the surface of the coating, the Zr-oxides consisted of ZrO2 and ZrSiO4 were gradually peeled into tiny oxide particles. Owing to the superhigh melting temperature of ZrO2 (2700 \\x0e C) and ZrSiO4 (2500 \\x0e C), the stripped tiny Zr-oxides have not been melted, but gradually form Zr-oxides-rings with the fluid of SiO2 glass layer, while the unpeeled large Zr-oxides inlayed in the  \\x0c', '394  X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  Fig. 13.  (a) Backscattered SEM micrograph of the surface of the SiC coating modified with ZrB2 after oxidation at 1773 K for 50 h; (b) magnification of (a); (c) Spot EDS analyses.  SiO2 glass layer as large obstacles. The refractory Zr-oxides-rings and the large unpeeled Zr-oxides-obstacles construct a kind of skeleton inlayed in the SiO2 glass layer, which enhances the stability of the glass layer at super-high temperature. As shown in Fig. 15(d), during the whole oxidation test, owing to the different thermal coefficient of expansion, some microcracks unavoidably generated. However, owing to the existence of skeleton of Zroxides, the compound Zr-Si-O glass layer demonstrates the ability of restricting the growth of microcracks, resulting in the occurrence of crack turning, thinning and pinning. The role of microcracks restricting is able to reduce the possibility of the formation of large and long microcracks that will provide path for oxygen to erode the carbon matrix.  2ZrB2 (s) þ 5O2 /2ZrO2 þ 2B2O3  SiC (s) þ 2O2 (g) / SiO2 (s) þ CO2 (g)  B2O3 (s) / B2O3 (l) / B2O3 (g)  ZrO2(s) þ SiO2(s)/ZrSiO4(s)  3.  Conclusions  (4)  (5)  (6)  (7)  In conclusion, the SiC coating modified with ZrB2 was prepared on the graphite substrates by liquid phase sintering technique. By the method of carbothermal reduction reaction, the fine ZrB2 powders whose particle size is in the range of 120e800 nm were prepared. During heat-treatment, the existence of Si powders and silica sol create a fluid condition for liquid phase sintering. After the modification of ZrB2, the ZrB2-SiC coating presents obvious advantages in contrast with pure SiC. The start mass loss temperature from 540 \\x0eC to 800 \\x0e C with the of the graphite substrate rises addition of ZrB2. The average mass loss rate of the fastest mass loss zone drops from \\x006 \\x02 10 \\x003 mg cm \\x002 s \\x001 to \\x002 \\x02 10 \\x003 mg cm \\x002 s \\x001.  (a) Cross sectional morphology of coating after oxidation; (b) magnification of  Fig. 14. (a).  \\x0c', 'X. Ren et al.  /  Journal of Alloys and Compounds 769 (2018) 387e396  395  Fig. 15. Sketch map of oxidation protection mechanisms of the SiC coating modified with ZrB2.  The range of the fastest weightless zone under the protection of SiC coating modified with ZrB2 is much shorter than the pure SiC coating. The fast weight gain temperature range (from 700 \\x0e C to 1240 \\x0eC) in the TG curve of the ZrB2 powders effectively compensate the weakness of the SiC coating in this temperature range, which is caused by the increased amount of the generated B2O3. After oxidation at 1773 K for 50 h, the movement of the SiO2 glass layer peeled the Zr-oxides into tiny oxide particles to form Zroxides-rings. The refractory large unpeeled Zr-oxides-obstacles and the Zr-oxides-rings construct skeleton inlayed in the SiO2 glass layer, which exhibits the ability of restricting the growth of microcracks, consuming the crack propagation energy and reducing the possibility of the formation of large and long micro cracks.  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},{
  "_id": 17,
  "PDF": "Application of a counter-current gaseous diffusion model to the oxidation of hafnium carbide at 1200 to 1530°C.pdf",
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P i e r r e   r e a c t i o n   c a n   o c c u r   w i t h i n   t h e   p o r o u s   o x i d e   a n d ,   a n a l o g o u s   t o   c a r b o n   o x i   d a t i o n ,   t h e   p o s i t i o n   w h e r e   i t   o c c u r s   m a y   b e   t e r m e d   a   \" f l a m e   f r o n t . \"   T h e   C O 2   g a s   g e n e r a t e d   a t   t h e   f l a m e   f r o n t   d i f f u s e s   b o t h   i n w a r d   t o w a r d   t h e   c a r b i d e / o x i d e   i n t e r f a c e   a n d   o u t w a r d   t o w a r d   t h e   0 2   a t m o s p h e r e .   A n   a n a l y s i s   o f   t h e   c o u n t e r c u r r e n t   d i f f u s i o n   o f   t h e   C O / C O 2   a n d   O 2 / C O 2   g a s   m i x t u r e s ,   a n d   t h e   p o s i t i o n   o f   t h e   f l a m e   f r o n t ,   f o r m   t h e   b a s i s   o f   t h e   c o u n t e r c u r r e n t   g a s e o u s   d i f f u s i o n   m o d e l   p r e s e n t e d   h e r e .   I n   p r e v i o u s   w o r k ,   1 3   t h e   f o r m a t i o n   o f   i n t e r c o n n e c t e d   p o r e s   d u r i n g   t h e   o x i d a t i o n   o f   h a f n i u m   c a r b i d e   i n   o x y g e n n i t r o g e n   g a s   m i x t u r e s   h a s   b e e n   d e s c r i b e d   a n d   t h e   p a r a b o l i c   o x i d a t i o n   r a t e   p a r a m e t e r s   h a v e   b e e n   p r e   s e n t e d .   A   c o u n t e r c u r r e n t   g a s e o u s   d i f f u s i o n   m o d e l   h a s   b e e n   d e v e l o p e d   a n d   i s   p r e s e n t e d   h e r e   t o   d e s c r i b e   t h e   o x i d a t i o n   o f   h a f n i u m   c a r b i d e .   P r e d i c t i o n s   o f   t h e   m o d e l   a r e   c o m p a r e d   w i t h   t h e   e x p e r i m e n t a l   r e s u l t s .   E X P E R I M E N T A L   C O N D I T I O N S   A N D   R E S U L T S   T h e   e x p e r i m e n t a l   p r o c e d u r e s   a n d   k i n e t i c   c a l c u l a t i o n s   f o r   o u r   p a s t   w o r k   o n   t h e   o x i d a t i o n   o f   h a f n i u m   c a r b i d e   h a v e   b e e n   p r e s e n t e d   e l s e w h e r e .   1 3   T h e r e f o r e ,   o n l y   a   b r i e f   d i s c u s s i o n   o f   t h e   e x p e r i m e n t a l   c o n d i t i o n s ,   p r o c e d u r e s ,   a n d   r e s u l t s   i s   p r e s e n t e d   h e r e .   H a f n i u m   c a r b i d e   w a s   o x i d i z e d   i n   t h r e e   d i f f e r e n t   o x y g e n n i t r o g e n   g a s   m i x t u r e s   i n   t h e   t e m p e r a t u r e   r a n g e   1 2 0 0 1 5 3 0 ~   T h e   o x y g e n   p a r t i a l   p r e s s   u r e s   i n   t h e   t h r e e   g a s   m i x t u r e s   w e r e   0 . 0 2 ,   0 . 2 1   ( a i r ) ,   a n d   1 . 0   ( p u r e   o x y g e n ) .   A l l   t h r e e   g a s   m i x t u r e s   w e r e   a t   a t m o s p h e r i c   p r e s s u r e .   T h e r m o g r a v i m e t r i c   a n a l y s i s   w a s   u s e d   t o   m e a s u r e   t h e   i n c r e a s e   i n   t h e   m a s s   o f   t h e   s a m p l e s   d u r i n g   o x i d a t i o n .   T h e   c h a n g e   i n   m a s s   s q u a r e d   w a s   f o u n d   t o   b e   l i n e a r   w i t h   r e s p e c t   t o   t i m e   a n d   t h e   s l o p e   i s   t h e   p a r a b o l i c   r a t e   c o n s t a n t ,   k p .   T h e   p a r a b o l i c   r a t e   c o n s t a n t   w a s   f o u n d   t o   b e   v i r t u a l l y   i n d e p e n d e n t   o f   t e m p e r a t u r e ,   w i t h   a v e r a g e   v a l u e s   o f   9 . 7   (cid:141)   1 0   8   g 2 / c m 4   s e c   w i t h   a n   o x y g e n   p a r t i a l   p r e s s u r e   o f   0 . 0 2 ,   5 . 0   (cid:141)   1 0   7   g 2 / c m 4   s e c   i n   a i r ,   a n d   2 . 0   x   1 0   6   g 2 / c m 4   s e c   i n   p u r e   o x y g e n .   E x a m i n a t i o n   o f   t h e   o x i d e   s c a l e s   u s i n g   S E M   a n d   S T E M   m i c r o s c o p y   r e v e a l e d   a   v e r y   p o r o u s   s c a l e   c o n t a i n i n g   i n t e r c o n n e c t e d   s u b m i c r o n   p o r e s .   T h e   s i z e s   o f   t h e   p o r e s   v a r i e d   c o n s i d e r a b l y   w i t h i n   t h e   s c a l e :   f r o m   0 . 0 1   m i c r o n s   i n   d i a m e t e r   a t   t h e   c a r b i d e   i n t e r f a c e   t o   0 . 3   m i c r o n s   i n   d i a m e t e r   i n   t h e   o u t e r   p a r t s   o f   t h e   s c a l e .   T H E O R Y   T h e   i n c o m p a t i b i l i t y   b e t w e e n   0 2   a n d   C O   g a s e s   a b o v e   1 2 0 0 ~   l e a d   t o   t w o   s e p a r a t e   r e g i o n s   w i t h i n   t h e   p o r o u s   o x i d e .   N e a r   t h e   c a r b i d e ,   t h e   g a s e o u s   s p e c i e s   p r e s e n t   a r e   C O ,   C O 2 ,   a n d   N   2 .   I n   t h e   o u t e r   r e g i o n   o f   t h e   s c a l e ,   t h e   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   1 1 1   g a s e o u s   s p e c i e s   p r e s e n t   a r e   0 2 ,   C 0 2 ,   a n d   N   2 .   B e t w e e n   t h e s e   t w o   r e g i o n s ,   a   \" f l a m e   f r o n t \"   e x i s t s   w i t h i n   t h e   p o r e s   a t   w h i c h   0 2   a n d   C O   r e a c t   t o   f o r m   C O 2 .   T h i s   r e a c t i o n   i s   a   s i n k   f o r   g a s   m o l e c u l e s ,   w h i c h   r e s u l t s   i n   n e t   g a s   f l o w ,   f r o m   b o t h   r e g i o n s ,   t o w a r d   t h e   f l a m e   f r o n t .   O n   e i t h e r   s i d e   o f   t h e   f l a m e   f r o n t ,   C O   a n d   C O 2   c o u n t e r d i f f u s e   a n d   C O 2   a n d   0 2   c o u n t e r d i f f u s e   a s   s h o w n   i n   F i g .   1 .   T h e   r e a c t i o n   a t   t h e   c a r b i d e / o x i d e   i n t e r f a c e   i s . \"   H f C   +   3 C O 2   =   H f O 2   +   4 C O   ( 1 )   w h i l e   t h e   o v e r a l l   r e a c t i o n   i s :   H f C   +   2 0 2   =   H f O   2   +   C O   2   ( 2 )   F i g u r e   1   s h o w s   t h e   f l u x e s   J i   a c r o s s   b o t h   r e g i o n s   i n   t e r m s   o f   a   g e n e r a l   u n i t   f l u x   %   d . )   X   L L   H f C   %   t   O   ~   o   c ~   V   1 . 0   E   o . 9   ( ~   0 . 8   0 . 7   0 . 6   t o   0 . 5   1 3 . .   0 . 4   . \" ~   0 . 3   1 3 _   P o r o u s   H I O   2   <   J o o = 4 a   ~   J c 0 2 = 3 a   +   I   0   o   +   I   (cid:12) 9   . . \\' ~   1   \\'   ~ ~ ~ .     N 2     L   _   J o z = 2 a   J c o 2 = l a   o   %   0   +   0   o   J o = 2 a   J c = l a   A i r   0 . 1   0   0   1   D i s t a n c e   f r o m   t h e   C a r b i d e / O x i d e   I n t e r f a c e   F i g .   1 ,   F l u x e s ,   r e a c t i o n s ,   a n d   p a r t i a l   p r e s s u r e s   o f   0 2 ,   C O ,   C O 2 ,   a n d   N   2   a c r o s s   a   p o r o u s   h a f n i a   s c a l e   d u r i n g   o x i d a t i o n   o f   H f C   i n   a i r   a t   1 4 0 0 ~   T h e   o v e r a l l   r e a c t i o n   i s   H f C   +   2 0 2   =   H f O 2   +   C O >   \\x0c', '1 1 2   H o l c o m b   a n d   S t .   P i e r r e   \" a \"   ( m o l e s / c m   2   s e c ) .   T h e   f l u x   o f   e a c h   s p e c i e s   i s   e x p r e s s e d   i n   m u l t i p l e s   o f   a   ( l a ,   2 a ,   e t c . ) .   T h e   n e t   f l u x   o f   o x y g e n   a t o m s   i n w a r d   i s   t w i c e   t h a t   o f   c a r b o n   a t o m s   o u t w a r d ,   a s   r e q u i r e d   b y   E q s .   ( 1 )   a n d   ( 2 ) .   A n   a p p r o x i m a t i o n   o f   t h e   S t e f a n M a x w e l l   e q u a t i o n   4   i s   u s e d   t o   c a l c u l a t e   t h e   g a s e o u s   c o n c e n t r a t i o n   ( C i )   p r o f i l e s   a c r o s s   b o t h   r e g i o n s   o f   t h e   p o r o u s   h a f n i a   s c a l e :   J i   =     , e f f ~ , ~ X )   \\' I \\'   Z   4 \"   ( 3 )   j = l   i n   w h i c h   D i e f f   i s   a n   e f f e c t i v e   d i f f u s i o n   c o e f f i c i e n t   o f   g a s   s p e c i e s   i ,   J i   i s   t h e   f l u x   o f   i   i n   t h e   r e g i o n   a n d   x   i s   t h e   d i s t a n c e   f r o m   t h e   c a r b i d e / o x i d e   i n t e r f a c e .   T h e   s o l u t i o n s   t o   E q .   ( 3 )   a t   a n y   p o s i t i o n   x   i n   t h e   c a r b i d e s i d e   o f   t h e   f l a m e   f r o n t   a r e :   C c o =   ( C ~ o 4 c )   e x p (   a ( x   l ) ~   k ,   C D C O e f f   J   +   4 c   ( 4 )   ,   f a ( x     l ) \\' ~   C c o 2   =   ( C c o ~   +   3 c ) e x p   ~ , ~ )     3 c   ( 5 )   (   ,   a ( l x ) ~   C N ~   = e x p   I n   C N 2   C D N 2 e f f , \\\\ ]   ( 6 )   T h e   s o l u t i o n s   t o   E q .   ( 3 )   o n   t h e   a i r s i d e   o f   t h e   f l a m e   f r o n t   a r e :   f a ( L     x ) \\' ~   C %   =   ( C ~ 2     2 c )   e x p   \\\\ / C ~ \" D o 2 e f f   j   +   2 c   ( 7 )   C c 0 2   =   ( C ~ %   + c ) e x p ( a ( L ~ x ) \\\\ ]   c   ( 8 )   \\\\   C L I c o 2 e f f   ~ /   (   a ( L x ) . ~   C N 2   =   e x p   I n   C ~ 2   +   C D N 2 e f f   \\\\ ]   ( 9 )   i n   w h i c h   C i   h a s   u n i t s   o f   m o l e s / c m   3 ,   l   i s   t h e   d i s t a n c e   f r o m   t h e   c a r b i d e / o x i d e   i n t e r f a c e   t o   t h e   f l a m e   f r o n t ,   L   i s   t h e   d i s t a n c e   f r o m   t h e   c a r b i d e / o x i d e   i n t e r   f a c e   t o   t h e   o u t e r   e d g e   o f   t h e   s c a l e ,   c   i s   t h e   m o l a r   c o n c e n t r a t i o n ,   C / *   i s   t h e   v a l u e   o f   C i   a t   x   e q u a l   t o / ,   a n d   C ?   i s   t h e   v a l u e   o f   C i   a t   x   e q u a l   t o   L .   U s i n g   E q s .   ( 4 9 ) ,   t h e   c o n c e n t r a t i o n s   o f   C O   a n d   C O 2   a t   t h e   c a r b i d e / o x i d e   i n t e r f a c e   a r e   e q u a l   t o :   C c o   =   _ 4 c e x p   I   a L   D % e f f / 2 c   ) \\\\ ]   I n \\\\ [   o   + 4 c   ( 1 0 )   c D c o e r r   D c o e f r   2 C 6 2     ~ c   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   1 1 3   a n d   {   ( 2 c ) \\\\ ]   }   C c o 2   =   ( C ~ o 2   +   c )   e x p   \\\\ [   D o 2 e f f   I n   2 c     ~   +   2 c   L D c % e f f     C %   x e x p \\\\ [   a L   ~ O 2 e ~ l n ( c o 2 C   c c ) l _ 3 c   c O c o 2 e f f   C O 2 e f f   \\\\   0 2   ( 1 1 )   E q u a t i o n s   ( 1 0 )   a n d   ( 1 1 )   a r e   u s e d   t o   c a l c u l a t e   t h e   g a s   c o m p o s i t i o n s   b y   f i n d i n g   t h e   v a l u e   o f   t h e   f l u x ,   a ,   f o r   w h i c h   t h e   r a t i o   C ~ o / C 3 o 2   a t   t h e   c a r b i d e / o x i d e   i n t e r f a c e   i s   t h e   s a m e   a s   t h e   e q u i l i b r i u m   c o n s t a n t   o f   E q .   ( 1 ) .   I n   E q s .   ( 1 0 )   a n d   ( l   1 ) ,   t h e   f l u x   a n d   s c a l e   t h i c k n e s s   a r e   a l w a y s   f o u n d   t o g e t h e r   a s   t h e   p r o d u c t   a L ,   w h i c h   h a s   u n i t s   o f   p e r m e a b i l i t y .   T h e   p r o d u c t   a L   i s   a   c o n s t a n t   a n d   i s   i n d e p e n d e n t   o f   t i m e .   T h e   r e l a t i o n s h i p   b e t w e e n   t h e   f l u x   a   a n d   t h e   t h i c k n e s s   o f   t h e   s c a l e   L   i s   a l s o   d e s c r i b e d   b y :   L   =   V   a d t   ( 1 2 )   w h e r e   V   i s   t h e   m o l a r   v o l u m e   ( c m 3 / g m o l e )   o f   t h e   p o r o u s   h a f n i a .   E q u a t i o n   ( 1 2 )   s u m s   u p   a l l   t h e   o x y g e n   u s e d   i n   c o n v e r t i n g   t h e   H f C   t o   H f O   2   a n d   c o n v e r t s   t h i s   v a l u e   i n t o   t h e   s c a l e   t h i c k n e s s ,   L .   C o m b i n i n g   E q .   ( 1 2 )   w i t h   t h e   f a c t   t h a t   t h e   p r o d u c t   a L   i s   a   c o n s t a n t ,   o n e   c a n   s h o w   t h a t   t h e   r a t i o   L 2 / t   i s   e q u a l   t o :   L 2 / t   =   2 V ( a L )   ( 1 3 )   T h i s   r a t i o ,   a   c o n s t a n t ,   i s   e q u i v a l e n t   t o   a   p a r a b o l i c   r a t e   c o n s t a n t ,   k p   ~   h a v i n g   u n i t s   o f   c m 2 / s e c .   T h e   c o n v e r s i o n   f r o m   k p   ~   t o   k p   ( g 2 / c m 4   s e c )   i s :   / 2 0 \\\\   2   o   k p   =   ~ ~ )   k ~   ( 1 4 )   T h e   f a c t o r   o f   2 0   a r i s e s   f r o m   t h e   f a c t   t h a t   t h e   c o n v e r s i o n   o f   o n e   m o l e   o f   H f C   t o   H f O 2   i n c r e a s e s   t h e   m a s s   b y   2 0   g r a m s .   T h e   a b o v e   e q u a t i o n s   u t i l i z e   a n   e f f e c t i v e   d i f f u s i o n   c o e f f i c i e n t   D i e f f .   F o r   d i f f u s i o n   i n   p o r o u s   o x i d e s ,   a n   e f f e c t i v e   d i f f u s i o n   c o e f f i c i e n t   i s   n e e d e d   w h i c h   d e c r e a s e s   t h e   o v e r a l l   f l u x   o f   g a s   d u e   t o   t h e   e f f e c t s   o f   p o r o s i t y ,   t o r t u o s i t y ,   a n d   g a s   m o l e c u l e   i n t e r a c t i o n s   w i t h   p o r e   w a l l s .   F o r   g a s   s p e c i e s   i ,   t h e s e   e f f e c t s   i n t r o d u c e   a n   e f f e c t i v e   m o l e c u l a r   d i f f u s i o n   c o e f f i c i e n t   D M i e f f   a n d   a n   e f f e c t i v e   K n u d s e n   d i f f u s i o n   c o e f f i c i e n t   D M i e f f ,   w h i c h   a r e   r e l a t e d   t o   D i ~ f f   b y :   4   1   D i e f f     ( 1     a N i ) / D M i e f f   +   1 / D K i e f f   ( 1 5 )   \\x0c', '1 1 4   H o i c o m b   a n d   S t .   P i e r r e   w h e r e   N i   i s   t h e   m o l e   f r a c t i o n   o f   i   a n d   a   i s   e q u a l   t o :   o ~   =   1     ( M i / M N 2 )   1 / 2   ( 1 6 )   a n d   w h e r e   M   i   i s   t h e   m o l e c u l a r   w e i g h t   o f   i .   T h e   e f f e c t i v e   K n u d s e n   d i f f u s i o n   c o e f f i c i e n t   c a n   b e   e x p r e s s e d   b y :   4   D , , , : i e f f   =   ( 4 / 3 ) p ,   i K   ( 1 7 )   w h e r e   # i   i s   t h e   m o l e c u l a r   v e l o c i t y   o f   i ,   a n d   K   i s   t h e   K n u d s e n   p e r m e a b i l i t y .   E v a n s ,   W a t s o n ,   a n d   M a s o n   5   h a v e   s h o w n   t h a t   f o r   a   c a p i l l a r y ,   K   i s   e q u a l   t o   o n e   h a l f   o f   t h e   a v e r a g e   p o r e   r a d i u s   r .   T h i s   i s   u s e d   a s   t h e   e s t i m a t e   o f   K   f o r   d e t e r m i n i n g   D x i ~ f f .   T h e   e f f e c t i v e   m o l e c u l a r   d i f f u s i o n   c o e f f i c i e n t   i s   u s u a l l y   e x p r e s s e d   a s :   D M i e f f   =   ( r   ( 1 8 )   w h e r e   r   i s   t h e   f r a c t i o n a l   p o r o s i t y ,   7   i s   t h e   t o r t u o s i t y ,   a n d   D i m   i s   a n   e f f e c t i v e   b i n a r y   d i f f u s i v i t y   i n   a   m u l t i c o m p o n e n t   s y s t e m .   4   K i m ,   O c h o a ,   a n d   W h i t a k e r   6   h a v e   d e t e r m i n e d   t h a t :   7   =   q ~ 0 . 4   ( 1 9 )   g i v e s   a   g o o d   e m p i r i c a l   f i t   f o r   r e l a t i n g   p o r o s i t y   t o   t o r t u o s i t y   f o r   p o r o s i t i e s   l e s s   t h a n   5 0 0 .   I t   h a s   b e e n   s h o w n   t h a t   D ~   c a n   b e   f o u n d   b y   c o m b i n i n g   E q .   ( 3 )   w i t h   t h e   S t e f a n M a x w e l l   e q u a t i o n :   O x   ( N i J j     N j J i )   ( 2 0 )   i n   w h i c h   t h e   b i n a r y   i n t e r d i f f u s i o n   c o e f f i c i e n t   D / j   i s   e x p r e s s e d   b y   u s i n g   t h e   L e n n a r d J o n e s   6 1 2   p o t e n t i a l :   3   T 3 / 2   ( 1   +   1 . \\\\ ] 1 / 2   D / j ( c m 2 / s e c )   =   1 . 8 5 8 3   (cid:141)   1 0   ~   M i   ~ \\\\ ]   ( 2 1 )   w h e r e   a   i s   t h e   c o l l i s i o n   d i a m e t e r   i n   a n g s t r o m s   a n d   f ~   i s   t h e   c o l l i s i o n   i n t e g r a l .   W h e n   E q s .   ( 1 5 ) ( 2 1 )   a r e   a p p l i e d   t o   t h e   c a l c u l a t i o n   o f   t h e   c o n c e n t r a t i o n   p r o f i l e s ,   a n   a v e r a g e   v a l u e   o f   N i   i s   u s e d   i n   E q .   ( 2 0 )   f o r   e a c h   g a s   s p e c i e   o n   b o t h   s i d e s   o f   t h e   f l a m e   f r o n t .   I t e r a t i v e   c a l c u l a t i o n s   a r e   u s e d   t o   e n s u r e   s e l f   c o n s i s t e n c y   b e t w e e n   t h e   r e l a t e d   c o n c e n t r a t i o n   p r o f i l e s   a n d   e f f e c t i v e   d i f f u s i o n   c o e f f i c i e n t s .   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   1 1 5   T a b l e   I .   C a l c u l a t e d   d i f f u s i o n   c o e f f i c i e n t s   a t   1 2 0 0   a n d   1 5 3 0 ~   f o r   a   p o r o s i t y   o f   0 . 0 2   a n d   a n   a v e r   a g e   p o r e   r a d i u s   o f   0 . 0 1   # m   S p e d e s   S i d e   o f   1 2 0 0 ~   1 5 3 0 ~   f l a m e   f r o n t   D / m   ( c m 2 / s e c )   D i e f f   ( c m 2 / s e c )   D i m   ( c m 2 / s e c )   D i e f f   ( c m 2 / s e c )   C O   C a r b i d e   2 . 9 0   0 . 0 1 0 1   4 . 0 5   0 . 0 1 3 6   C O 2   C a r b i d e   2 . 3 3   0 . 0 0 8 0   3 . 2 5   0 . 0 1 0 7   N a   C a r b i d e   2 . 7 1   0 . 0 0 9 6   3 . 7 9   0 . 0 1 2 9   0 2   A i r   2 . 9 2   0 . 0 1 0 0   4 . 0 8   0 . 0 1 3 4   C O 2   A i r   2 . 3 3   0 . 0 0 8 0   3 . 2 5   0 . 0 1 0 7   N   2   A i r   2 . 8 0   0 . 0 0 9 9   3 . 9 1   0 . 0 1 3 2   T h e   d i f f u s i o n   c o e f f i c i e n t s   f o u n d   i n   t h e s e   c a l c u l a t i o n s   a r e   g i v e n   i n   T a b l e   I .   T h e   e f f e c t s   o f   p o r o s i t y ,   t o r t u o s i t y ,   a n d   i n t e r a c t i o n s   w i t h   p o r e   w a l l s   d e c r e a s e   t h e   d i f f u s i o n   c o e f f i c i e n t   b y   a   f a c t o r   o f   3 0 0 .   F i g u r e   1   s h o w s   t h e   c a l c u l a t e d   g a s   c o n c e n t r a t i o n   p r o f i l e s ,   f r o m   E q s .   ( 4 ) ( 9 ) ,   i n   t e r m s   o f   p a r t i a l   p r e s s u r e s ,   a c r o s s   b o t h   s i d e s   o f   t h e   f l a m e   f r o n t .   C O M P A R I S O N   O F   T H E   M O D E L   W I T H   E X P E R I M E N T A L   R E S U L T S   T h e   r e l a t i o n s h i p   b e t w e e n   p o r o s i t y ,   a v e r a g e   p o r e   r a d i u s ,   a n d   t h e   p a r a b o l i c   r a t e   c o n s t a n t   i s   i l l u s t r a t e d   i n   F i g .   2 ,   w h i c h   s h o w s   a   f a m i l y   o f   c u r v e s   h a v i n g   d i f f e r e n t   p a r a b o l i c   r a t e   c o n s t a n t s .   T h e   c u r v a t u r e   i s   a   r e s u l t   o f   E q .   ( 1 5 ) ,   a n d   s e p a r a t e s   t h e   g r a p h   i n t o   t h r e e   r e g i m e s :   w h e r e   K n u d s e n   t y p e   d i f f u s i o n   d o m i n a t e s ,   w h e r e   v i s c o u s   f l o w   d o m i n a t e s ,   a n d   a   m i x e d   r e g i m e   w h e r e   b o t h   c o n t r i b u t e .   3 0 0   ~   v   \\\\ [   k p   =   6 . 8 X 1 0   \" 7   ( g 2 / c m 4 s e c )   U )   /         k p   =   5 . 4 x 1 0   . 7   ( g 2 / c m 4 s e c )   . ~ _ _ .   I     k p   =   4 . 1 x 1 0   \" 7   ( g 2 / c m 4 s e c )   2 0 0   c C   o   ~   0 . .   1 0 0   >   <   0   I   ,   I   ~   I   ~   I   0 . 0 1   0 . 0 2   0 . 0 3   0 . 0 4   0 . 0 5   P o r o s i t y   F i g .   2 .   T h e   p a r a b o l i c   r a t e   c o n s t a n t   f o r   H f C   o x i d a t i o n   i n   a i r   a t   1 4 0 0 ~   a s   a   f u n c t i o n   o f   p o r o s i t y   a n d   a v e r a g e   p o r e   r a d i u s .   \\x0c', \"1 1 6   H o l c o m b   a n d   S t .   P i e r r e   T h e   h a f n i a   s c a l e s   h a v e   c o m p l e x   s t r u c t u r e s   c o n t a i n i n g   a   r a n g e   o f   p o r o s i t i e s   a n d   p o r e   r a d i i .   N e a r   t h e   c a r b i d e / o x i d e   i n t e r f a c e   t h e   p o r o s i t y   a n d   p o r e   s i z e   h a v e   b e e n   o b s e r v e d   t o   b e   m u c h   s m a l l e r   t h a n   i n   t h e   r e s t   o f   t h e   s c a l e .   O n e   c a n n o t   e x p e c t   t o   d e s c r i b e   t h e   e n t i r e   s c a l e   u s i n g   j u s t   t h e s e   t w o   p a r a m e t e r s   o f   p o r o s i t y   a n d   a v e r a g e   p o r e   r a d i u s .   O n e   c a n   h o w e v e r   u s e   t h e s e   t w o   p a r a m e t e r s   t o   d e s c r i b e   a   s c a l e   t h a t   h a s   e q u i v a l e n t   t r a n s p o r t   p r o p e r   t i e s .   A s   F i g .   2   s h o w s ,   m a n y   c o m b i n a t i o n s   o f   p o r o s i t y   a n d   a v e r a g e   p o r e   r a d i u s   r e s u l t   i n   t h e   s a m e   o v e r a l l   t r a n s p o r t   p r o p e r t i e s .   F i g u r e   3   s h o w s   a   c o m p a r i s o n   i n   p a r a b o l i c   r a t e   c o n s t a n t s   b e t w e e n   e x p e r   i m e n t a l   r e s u l t s   1 3   a n d   t h e   m o d e l   f o r   a   s c a l e   w i t h   a   p o r o s i t y   o f   0 . 0 2   a n d   a n   a v e r a g e   p o r e   r a d i u s   o f   0 . 0 1   # m .   T h i s   c o m b i n a t i o n   o f   p a r a m e t e r s   l i e   i n   t h e   m i x e d   r e g i o n   o f   F i g .   2 .   T h e   p a r a m e t e r s   q ~   a n d   r   a r e   m u c h   s m a l l e r   t h a n   w h a t   w o u l d   b e   f o u n d   b y   a v e r a g i n g   t h e   v a l u e s   o f   ~ b   a n d   r   t h r o u g h o u t   t h e   5 . 0   o   o   o   ( . 5   E : ,   o   b   o   O   O   O   O   t O   ' ~   C O   ~ 1   I   I   I   I   5 . 5   P O   2   =   1 . 0 0   O   P O   2   =   0 . 2 1   (cid:12) 9   (cid:12) 9   . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .   \\\\ [ \\\\ ]   P O   2   =   0 . 0 2   \\\\ [ \\\\ ]   \\\\ [ \\\\ ]   . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .   7 . 5   O   P O   2   =   1 . 0 0   (cid:12) 9   P O   2   =   0 . 2 1   P o r o s i t y   =   0 . 0 2   P O   2   =   0 . 0 2   A v e r a g e   P o r e   R a d i u s   =   1 0 0 , ~   6 . 0   _ l   \\\\ [   q   I   ~   ~   i   \\\\ ]   i   I   i   I   i   I   i   5 . 4   5 . 6   5 . 8   6 . 0   6 . 2   6 . 4   6 . 6   6 . 8   7 . 0   1 0 0 0 0 / T e m p e r a t u r e   ( K )   F i g .   3 .   C o m p a r i s o n   o f   e x p e r i m e n t a l   H f C   o x i d a t i o n   r e s u l t s   1 3   w i t h   t h e   c o u n t e r c u r r e n t   g a s e o u s   d i f f u s i o n   m o d e l   a s   a   f u n c t i o n   o f   r e c i p r o c a l   t e m p e r a t u r e .   o   6 . O   ( D   0 O   E   O   ~   6 . 5   . 0 0   7 . 0   \\x0c\", 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   1 1 7   o b s e r v e d   h a f n i a   s c a l e s .   H o w e v e r ,   t h e y   a p p e a r   t o   b e   c l o s e   i n   s i z e   t o   t h o s e   f o u n d   n e a r   t h e   c a r b i d e / o x i d e   i n t e r f a c e .   1 3   T h e   c o m p a r i s o n   b e t w e e n   t h e   e x p e r i   m e n t a l   r e s u l t s   a n d   t h e   m o d e l   i s   a   f a v o r a b l e   o n e   i n   t e r m s   o f   t h e   e f f e c t s   o f   o x y g e n   p a r t i a l   p r e s s u r e   a n d   t e m p e r a t u r e .   T h i s   s u g g e s t s   t h a t   t h e   o x i d a t i o n   k i n e t i c s   a r e   l i m i t e d   b y   t h e   r e l a t i v e l y   d e n s e   s c a l e   n e a r   t h e   c a r b i d e / o x i d e   i n t e r f a c e .   A P P L I C A T I O N   O F   T H E   M O D E L   T O   O T H E R   S Y S T E M S   T h e   c o u n t e r c u r r e n t   g a s e o u s   d i f f u s i o n   m o d e l   w a s   d e v e l o p e d   f o r   t h e   o x i d a t i o n   o f   H f C .   A d d i t i o n a l   a p p l i c a t i o n s   o f   t h i s   m o d e l   c e r t a i n l y   i n c l u d e   t h e   o x i d a t i o n   o f   o t h e r   m e t a l   c a r b i d e   s y s t e m s ,   s u c h   a s   Z r C ,   w h i c h   f o r m   v i s c o u s   o x i d e s .   T h e   m o d e l   i s   a l s o   a p p l i c a b l e   t o   h i g h   t e m p e r a t u r e   c o a t i n g   s y s t e m s   i n   w h i c h   a   p o r o u s   o x i d e   o v e r c o a t   i s   a p p l i e d   t o   h e l p   r e s i s t   e r o s i o n   o x i d a t i o n ,   t o   i m m o b i l i z e   a   p r o t e c t i v e   g l a s s   f i l m ,   o r   t o   s e r v e   a s   a   t h i c k   s t a g n a n t   b o u n d a r y   l a y e r .   T h e   r a t e   o f   a t t a c k   t h r o u g h   t h e   p o r o u s   o x i d e   o v e r c o a t   c o u l d   b e   d e s c r i b e d   w i t h   t h e   m o d e l .   T h e   m o d e l   h a s   a l s o   b e e n   u s e d   t o   d e s c r i b e   t h e   o x i d a t i o n   o f   c a r b o n   c a r b o n   c o m p o s i t e s   t h r o u g h   a   c r a c k e d   S i C   c o a t i n g   d u r i n g   c o o l i n g   f r o m   e l e v a t e d   t e m p e r a t u r e s .   1   I n   t h i s   a p p l i c a t i o n ,   t h e   v a l u e   o f   t h e   e q u i l i b r i u m   c o n s t a n t   b e t w e e n   C ,   C O ,   a n d   C O 2   c h a n g e s   r a p i d l y   b e l o w   8 0 0 ~   w h i c h   r e s u l t s   i n   a   p r e d i c t e d   s h i f t   o f   t h e   f l a m e   f r o n t   t o w a r d   t h e   c a r b o n c a r b o n   c o m   p o s i t e .   T h i s   s h i f t   i n   t h e   p o s i t i o n   o f   t h e   f l a m e   f r o n t   r e d u c e s   t h e   c o n c e n t r a t i o n   g r a d i e n t   o f   i n c o m i n g   0 2 ,   w h i c h   l o w e r s   t h e   r a t e   o f   o x i d a t i o n   b e l o w   8 0 0 ~   I n   a n   i n d e p e n d e n t   i n v e s t i g a t i o n ,   B e r n s t e i n   a n d   K o g e r   7   d e v e l o p e d   a   s i m i l a r   m o d e l   f o r   c a r b o n   f i l m   u n d e r c u t   k i n e t i c s   i n   p u r e   o x y g e n .   T h e   k i n e t i c s   o f   a   p r o c e s s   w a s   p r e s e n t e d   f o r   f a b r i c a t i n g   m i c r o m e c h a n i c a l   s t r u c t u r e s   i n   w h i c h   a   s a c r i f i c i a l   l a y e r   o f   c a r b o n   i s   d e p o s i t e d   o n   a   s u b s t r a t e ,   f o l l o w e d   b y   a   t o p   l a y e r   o f   a   d i f f e r e n t   m a t e r i a l .   A f t e r   o x i d a t i o n   o f   t h e   c a r b o n   l a y e r ,   t h e   t o p   l a y e r   i s   l e f t   f r e e .   O n e   o f   t h e   b o u n d a r y   c o n d i t i o n s   u s e d   w a s   p u r e   C O   g a s   a t   t h e   c a r b o n   i n t e r f a c e ,   r a t h e r   t h a n   u s i n g   a n   e q u i l i b r i u m   c o n s t a n t   t o   r e l a t e   t h e   C O   a n d   C O 2   p a r t i a l   p r e s s u r e s   a s   w a s   d e s c r i b e d   a b o v e   f o r   t h e   o x i d a t i o n   o f   c a r b o n c a r b o n   c o m p o s i t e s   t h r o u g h   c r a c k e d   S i C   c o a t i n g s .   T h i s   r e s u l t s   i n   n o t   p r e d i c t i n g   a   s h i f t   i n   t h e   p o s i t i o n   o f   t h e   f l a m e   f r o n t   t o w a r d   t h e   c a r b o n   i n t e r f a c e   a n d   t h e r e f o r e   n o   s t r o n g   t e m p e r a t u r e   d e p e n   d e n c e .   Y e t   b e l o w   7 0 0 ~   t h e   m e a s u r e d   o x i d a t i o n   r a t e s   w e r e   f o u n d   t o   b e   e x t r e m e l y   s l o w ,   a s   w o u l d   b e   p r e d i c t e d   b y   a   s h i f t   i n   t h e   f l a m e   f r o n t .   S U M M A R Y   A N D   C O N C L U S I O N S   A   c o u n t e r c u r r e n t   g a s e o u s   d i f f u s i o n   m o d e l   w a s   d e v e l o p e d   f o r   d e s c r i b i n g   t h e   o x i d a t i o n   o f   H f C   t h r o u g h   p o r o u s   h a f n i a   b e t w e e n   1 2 0 0   a n d   1 5 3 0 ~   T h e   \\x0c', '1 1 8   H o l e o m b   a n d   S t .   P i e r r e   m o d e l   s e p a r a t e s   t h e   p o r o u s   h a f n i a   i n t o   t w o   g a s   d i f f u s i o n   r e g i o n s   s e p a r a t e d   b y   a   f l a m e   f r o n t ,   w h e r e   0 2   a n d   C O   r e a c t   t o   f o r m   C O 2 .   I n   t h e   o u t e r   r e g i o n   ( a i r   s i d e )   0 2   a n d   C O 2   c o u n t e r d i f f u s e ;   i n   t h e   i n n e r   r e g i o n   ( c a r b i d e   s i d e )   C O 2   a n d   C O   c o u n t e r d i f f u s e .   T h e   s i m p l i f i e d   f o r m   o f   t h e   S t e f a n M a x w e l l   e q u a t i o n ,   c o u p l e d   w i t h   t h e   e f f e c t s   o f   p o r o s i t y ,   t o r t u o s i t y ,   a n d   g a s w a l l   i n t e r a c t i o n s   a r e   i n t r o d u c e d .   T h e   c o n c e n t r a t i o n   g r a d i e n t s   o f   e a c h   g a s e o u s   s p e c i e   i n   t h e   p o r e s   o f   t h e   h a f n i a   a r e   d e t e r m i n e d   a n d   t h e   r a t e   o f   o x i d a t i o n   i s   c a l c u l a t e d .   I t   w a s   s h o w n   t h a t   t h e   o x i d a t i o n   k i n e t i c s   p r e d i c t e d   b y   t h e   m o d e l   a r e   p a r a b o l i c .   A   p o r o s i t y   o f   2 %   a n d   a   p o r e   r a d i u s   o f   0 . 0 1   # m   a r e   r e p r e s e n t a t i v e   o f   t h e   v a l u e s   o b s e r v e d   i n   h a f n i a   d u r i n g   t h e   e a r l y   s t a g e s   o f   H f C   o x i d a t i o n .   T h e s e   v a l u e s   l e a d   t o   p r e d i c t i o n s   o f   p a r a b o l i c   r a t e   c o n s t a n t s   t h a t   a r e   c l o s e   t o   t h o s e   m e a s u r e d   b y   t h e r m o g r a v i m e t r i c   a n a l y s i s .   I n   a d d i t i o n ,   t h e   p r e d i c t e d   a n d   m e a s u r e d   p a r a b o l i c   r a t e   c o n s t a n t s   a r e   s h o w n   t o   h a v e   t h e   s a m e   d e p e n d e n c e   u p o n   t e m p e r a t u r e   a n d   o x y g e n   p a r t i a l   p r e s s u r e .   A p p l i c a t i o n s   o f   t h e   m o d e l   t o   o t h e r   s y s t e m s   a r e   p o s s i b l e .   O t h e r   a p p l i   c a t i o n s   i n c l u d e   t h e   o x i d a t i o n   o f   o t h e r   m e t a l   c a r b i d e s   s u c h   a s   Z r C ,   t h e   u s e   o f   p o r o u s   o x i d e   o v e r c o a t s   i n   h i g h t e m p e r a t u r e   c o a t i n g   s y s t e m s ,   a n d   t h e   o x i d a t i o n   o f   c a r b o n   ( o r   c a r b o n c a r b o n   c o m p o s i t e s )   t h r o u g h   c r a c k s .   A C K N O W L E D G M E N T S   S u p p o r t   o f   t h e   o f f i c e   o f   N a v a l   R e s e a r c h   a n d   t h e   U . S .   A i r   F o r c e   i s   a p p r e c i a t e d   b y   t h e   a u t h o r s .   R E F E R E N C E S   1 .   G .   R .   H o l c o m b ,   P h . D .   t h e s i s ,   T h e   O h i o   S t a t e   U n i v e r s i t y   ( 1 9 8 8 ) .   2 .   J . T .   P r a t e r ,   M o d i f i c a t i o n   o f   H a f n i u m   C a r b i d e   f o r   E n h a n c e d   O x i d a t i o n   R e s i s t a n c e   T h r o u g h   A d d i t i o n s   o f   T a n t a l u m   a n d   P r a s e o d y m i u m ,   A i r   F o r c e   R e p o r t   A F W A L T R 8 8 4 1 4 1 ,   1 9 8 8 .   3 .   E .   L .   C o u r t r i g h t ,   J .   T .   P r a t e r ,   G .   R .   H o l c o m b ,   G .   R .   S t .   P i e r r e ,   a n d   R .   A .   R a p p ,   O x i d .   M e t .   3 6 ,   4 2 3 4 3 7   ( 1 9 9 1 ) .   4 .   K .   S c h w e r d t f e g e r   a n d   E .   T .   T u r k d o g a n ,   i n   P h y s i c o c h e m i c a l   M e a s u r e m e n t s   i n   M e t a l s   R e s e a r c h   ( V o l .   4 ,   P a r t   I ,   C h a p .   4 B ) ,   R .   A .   R a p p ,   e d .   ( 1 9 7 0 ) ,   p p .   3 2 1 4 0 7 .   5 .   R .   B .   E v a n s ,   G .   M .   W a t s o n ,   a n d   E .   A .   M a s o n ,   J .   C h e m .   P h y s .   3 5 ,   2 0 7 6   ( 1 9 6 1 ) .   6 .   J .   H .   K i m ,   J .   A .   O c h o a ,   a n d   S .   W h i t a k e r ,   T r a n s .   P o r o u s   M e d i a   2 ,   3 2 7 3 5 6   ( 1 9 8 7 ) .   7 .   J .   B e r n s t e i n   a n d   T .   B .   K o g e r ,   J .   E l e c t r o c h e m .   S o c .   8 ,   2 0 8 6 2 0 9 0   ( 1 9 8 8 ) .   \\x0c']"
},{
  "_id": 18,
  "PDF": "Arc-jet testing on HfB2 and HfC-based ultra-high temperature ceramic materials.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 28 (2008) 1899-1907  Arc-jet testing on HfB2 and HfC-based ultra-high temperature ceramic materials  Raffaele Savino a,∗  , Mario De Stefano Fumo a , Laura Silvestroni b , Diletta Sciti b  a Dipartimento di Ingegneria Aerospaziale, University of Naples “Federico II”, P.le V. Tecchio 80, 80125 Naples, Italy b ISTEC, Institute of Science and Technology for Ceramics, CNR, Via Granarolo 64, 48018 Faenza, Italy  Received 25 July 2007; received in revised form 13 November 2007; accepted 25 November 2007  Available online 4 March 2008  Abstract  The behaviour of pressureless sintered HfC and HfB2 ceramics, when exposed to high enthalpy plasma flows typical of atmospheric re-entry environment, was investigated with an arc-jet facility at temperatures exceeding 2000 C. The surface temperature and emissivity of the materials were evaluated during the test. The microstructure modifications were analysed after exposure. Fluid dynamic numerical simulations were carried out to evaluate the catalytic atom recombination efficiencies of the materials at the experimental conditions. Surface and cross sections of the samples showed the formation of scales mainly consisting of HfO2 and SiO2 . For the HfB2 -based composite numerical results correlated quite well with experimental ones assuming a low catalytic surface behaviour. For the HfC-based material the surface behaviour changed from low catalytic to partially catalytic as the temperature increased. The post-test analyses confirm the potential of these composites to endure re-entry conditions with temperature approaching 2000 C or even higher. © 2008 Elsevier Ltd. All rights reserved.        Keywords: Arc-jet; Thermal protection systems; HfB2 ; HfC; Oxidation  1.  Introduction     Ultra-high temperature ceramics (UHTCs) are currently considered as emerging materials for aerospace applications.1-4 The increasing attention is driven by the demand of developing reusable hot structures as thermal protection systems (TPS) of re-entry vehicles characterised by sharp leading edges and therefore by larger aerothermal heating than blunt edges, such as those on the Space Shuttle, able to withstand temperatures that may exceed 2000 C during re-entry. As available materials cannot survive such extreme temperatures, new ones are required for advanced thermal protection systems.1,4,5 The use of UHTCs for sharp leading edges would also imply lower aerodynamic drag, improved flight performances and crew safety, due to the larger cross range and manoeuvrability along with more gentle re-entry trajectories.3,6,7 Hafnium boride and hafnium carbide, belonging to the class of the UHTCs, are candidates for thermal protection materials  ∗  Corresponding author. Tel.: +39 0817682357; fax: +39 0815932044.  E-mail address: rasavino@unina.it (R. Savino).  0955-2219/$ - see front matter © 2008 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2007.11.021        in both re-entry and hypersonic vehicles because of their high melting points (3900 C) and excellent chemical stability.8-13 Other notable properties are their high hardness, high electrical and thermal conductivity.8-13 Despite all the potentialities, so far these compounds have not been developed on industrial scale due to the difficult sinterability and low fracture toughness. Recent studies have pointed out that the addition of MoSi2 as sintering aid allows the achievement of highly C by pressureless sintering.14,15 dense bodies (98%) at 1950 Furthermore, the addition of MoSi2 is expected to improve the oxidation resistance due to the development of a silica protective coating.16 In this paper, arc-jet testing at temperatures between 1950 C and 2400 C is carried out on pressureless sintered HfB2 and HfC-based materials. Arc-jet testing represents the best groundbased simulation of a re-entry environment, in different ways. On one hand, it provides the possibility to explore the oxidation behaviour of these materials under extreme conditions. On the other hand, the materials response to large heat fluxes is evaluated through the determination of two important parameters, i.e. emissivity and catalytic efficiency. High values of emissivity and low values of surface catalytic efficiency are desired for        \\x0c', '1900  Table 1  R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  Starting materials: compositions, densities, thermal properties and emissivity  Label  Composition (vol%)  Sintering cycle  Bulk density (g/cm3 )  Relative  density (%)  Mean grain size (\\u242em)  HB5  HC5  HfB2 + 5% MoSi2 HfC + 5% MoSi2       1950  1950  C/60 min  C/60 min  10.7  12.1  98  98  1.5  3.0  Emissivity  0.9 (1600-2000  0.7 (1800-2400       C)  C)  2. Experimental  2.1. Material processing and characterisation  The following materials were selected for the arc-jet tests:  HfB2 + 5 vol% MoSi2 , labelled as HB5. HfC + 5 vol% MoSi2 , labelled as HC5.  Commercial powders were used to prepare the ceramic materials: HfB2 (Cerac Incorporated, Milwaukee, USA), particle size range 0.5-5 \\u242em, impurities: Al (0.07%), Fe (0.01%), Zr (0.47%); Cubic HfC (Cerac Inc., USA), 325 mesh, fisher size 1.1 \\u242em, grain size range 0.2-1.5 \\u242em; Tetragonal MoSi2 (<2 \\u242em, Aldrich, USA), mean particle size 2.8 \\u242em, grain size range 0.3-5 \\u242em and oxygen content 1 wt%. The powder mixtures were ultrasonically treated and milled for 24 h in absolute ethanol using zirconia milling media, then dried in a rotary evaporator and sieved to −250 mesh screen size. Four-centimetre diameter pellets were linearly pressed and subsequently cold isostatically pressed under 350 MPa before sintering. The pellets were pressureless sintered in a resistanceheated graphite furnace under a flowing argon atmosphere (1 atm) at 1950 C for 60 min. The bulk density was measured with the Archimedes method. The relative density was calculated dividing the bulk density by the theoretical density that was evaluated with the rule of mixture on the basis of the starting compositions. The dense samples were examined using X-ray diffraction (Siemens D500, Germany) to identify crystalline phases. The microstructures were polished with diamond to 0.25 \\u242em and were paste analysed with scanning electron microscopy (SEM, Cambridge S360) and energy dispersive spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK). The main properties of the processed materials are summarized in Table 1. After the arc-jet tests, the ceramic models     Fig. 1. Hemispheric HfB2 model used for R = 7.5 mm.  arc-jet  testing,  curvature  radius  the above mentioned applications as they reduce temperature gradients and thermal stresses in the structure, thus enabling the vehicle to operate under relatively high enthalpy flow conditions. So far, ZrB2 -SiC and HfB2-SiC composites were the predecessor materials analysed by arc-jet testing in the literature of UHTCs.3,17 Microstructural modifications induced by high thermal loading are investigated and discussed. In addition, fluid dynamic numerical simulations are carried out in order to rebuild, through computational fluid dynamic (CFD) modelling, the experimental tests and to evaluate an average catalytic efficiency of the different materials with respect to oxygen and nitrogen surface recombination reactions.  Fig. 2. Experimental setup. (a) Before the test and (b) during the test.  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  1901  were further analysed by SEM-EDS on surface and cross section.  2.2. Plasma torch tests  Samples with a hemispheric shape (Fig. 1) were machined through diamond-loaded tools and then exposed to sustained enthalpy flows using the arc-jet facility equipped with a 80 kW plasma torch that operates in inert gas (He, N2 , Ar and their mixtures) at mass flow rates up to 5 g/s. The specimens were located at a distance of 6 cm from the exit torch. (Fig. 2a and b) The HfB2 -based model was tested with an initial average specific total enthalpy of about 20 MJ/kg, that was gradually increased tuning the arc current up to 26 MJ/kg and then maintained for approximately 30 s. HfC-based models were exposed to hot streams at two different conditions. The first test denoted as HC5-I, was conducted setting the initial average specific total enthalpy at 20 MJ/kg for about 40 s; then a value of 22 MJ/kg was achieved and maintained for about 60 s. During the second test, denoted as HC5-II, the specific total enthalpy was set at 22 MJ/kg during the first 90 s, 24 MJ/kg during the following 60 s and 26 MJ/kg during the last 240 s. During the experiments, infrared and optical windows in the test chamber allowed visual inspection and diagnostic analyses. An automatic control system monitored the main parameters of the apparatus (voltage and current of the arc heater, water cooling temperature and mass flow rate). In particular, the specific total enthalpy (H) was evaluated through an energy balance between the energy supplied to the gas by the arc heater and the energy transferred to the cooling system (measured by the water temperature jump between inlet and outlet). The output data, processed via a dedicated software, allowed the evaluation of the surface temperature profile versus exposure time of the mode. Due to the extremely high thermal loading upon the ceramic models, surface chemical reactions like oxidation can be responsible for changes in the material’s emissivity. To overcome this problem, the experiments’ measurements were carried out with a radiation ratio pyrometer (Infratherm ISQ5, Impac Electronic Gmbh, Germany) which operates both in two colour and in the single colour function. In the two colour mode the instrument makes use of the ratio of two spectral radiances, measured at different wavelengths (0.9-1.05 \\u242em), to evaluate the true temperature.  This overcomes the problem of the emissivity knowledge since it is supposed to be the same at both wavelengths. Once the temperature was measured with the ratio pyrometer, its value was input to evaluate the spectral emissivity using the single colour function (λ = 0.9 \\u242em). In combination with the pyrometer, an infrared thermo-camera (Thermacam SC 3000, FLIR Systems, USA) was used to measure the surface temperature distributions and the spectral emissivity in the long wave range of the thermograph (λ = 9 \\u242em).  2.3. Numerical simulation of the plasma torch flow  To assess if the environment generated by the plasma torch at atmospheric pressure was able to reproduce heat fluxes, temperatures and reactive environments typical of atmospheric re-entry conditions, a systematic numerical analysis was carried out. The computations were carried out solving the full Navier-Stokes equations for a turbulent multi-reacting gas mixture with six chemical species (Ar, O, O2 , NO, N and N2 ) in chemical nonequilibrium. Each species of the mixture was assumed to behave as a thermally perfect gas, where translational-rotational and vibrational-electronic degrees of freedom were characterised by different temperatures. Vibrational-translational energy exchanges were modelled according to the Landau-Teller model, while the vibrational relaxation time was derived from the Millikan-White formula, with Park correction for high temperatures.18 Chemical and vibrational non-equilibrium was taken into account using the Park model.19-21 The fluid dynamics equations were numerically solved in the computational domain (plasma torch and test chamber). Convective fluxes were computed according to Roe’s Flux Difference Splitting scheme. Integration of the equations was implicit in time performed, until steady state was achieved, solving the linearised system of equation by the multigrid technique.  3. Results and discussion  3.1. Microstructural features of the as-sintered samples  3.1.1. HfB2 -based composite  According to X-ray diffraction analysis (not shown), sintered specimen was constituted by hexagonal HfB2  the and  Fig. 3. Polished surfaces of (a) HfB2 and (b) HfC-based composites.  \\x0c', '1902  R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  this material. HfC grains had a squared shape and displayed a bright colour, while the MoSi2 phase had an irregular shape and was recognizable as a darker contrast phase. The mean grain size of the carbide grains was about 3 \\u242em (Table 1). Further details are reported in an earlier work.15  3.2. Plasma flow characterisation  The results herein presented refer to a 75% argon-25% nitrogen mixture plasma jet with mass flow rate of 1.45 g/s, for an average specific total enthalpy of the flow at the torch exit varying from 20 MJ/kg to 28 MJ/kg, at atmospheric pressure. At the exit of the torch the plasma containing argon, nitrogen and atomic nitrogen expands through a nozzle (5 mm in diameter), comes into contact with the surrounding air at ambient conditions, so that oxygen in the atmosphere dissociates and a reacting mixture composed of Ar, O2 , N2 , NO, O and N is formed. Fig. 4 shows the results of computations performed for the case of an average specific total enthalpy of 26 MJ/kg. According to the calculations, the average specific total enthalpy in proximity of the specimen reduces drastically to about 6-8 MJ/kg (Fig. 4). Table 2 summarizes the test conditions and the computed flow characteristics for the different experiments. Correspondingly, Fig. 5a and b shows the increase of the surface temperature as a function of the exposition time for HfB2 and HfC samples, respectively. It must be mentioned that the temperature reached at the sample surface depends on the ability of the material to reject the heat by radiation, i.e. on its emissivity (ε = 1 for an ideal black body, ε < 1 for a real material surface). The higher is the emissivity, the greater is the emitted radiation. However, the temperature of the sample also depends on its thermal conductivity, since a high thermal conductivity allows heat to be conducted from the leading edge to colder zones and from there to be radiated away. The maximum temperature reached on the surface of HfB2 sample was 1950 C, Fig. 5a. The corresponding stagnation point heat flux, computed by numerical simulation was in the range 5-8 MW/m2 . A value of about 0.9 was estimated     Fig. 4. Computed mass fractions of the different species along the torch axis  and specific total enthalpy contour. H = 26 MJ/kg.  tetragonal MoSi2 and monoclinic HfO2 .14 The relative density was 98%, as reported in Table 1. The polished section (Fig. 3a) showed a regular microstructure, with little residual porosity. HfB2 grains had a rounded shape with mean grain size estimated by image analysis of about 1.5 \\u242em, while the MoSi2 phase had an irregular morphology with very low dihedral angles (20-30 ). This peculiar characteristic indicates that MoSi2 was very ductile at the sintering temperature and was accommodated between the voids left by the HfB2 skeleton. Further details are reported elsewhere.14     3.1.2. HfC-based composite  Cubic HfC and tetragonal MoSi2 were the crystalline phases detected after sintering. The final density was 98% (Table 1). The typical microstructure of the HfC-based composite is shown in Fig. 3b. Very few porosity was detected by SEM inspections on  Table 2  Test conditions  Flow conditions  Exit torch conditions        Average specific total enthalpy (MJ/kg)  Temperature (  C)  Flow conditions at model location  Specific total enthalpy (MJ/kg)  Temperature (  C)  Ar mass fraction  N2 mass fraction N mass fraction  O2 mass fraction O mass fraction  NO mass fraction  Stagnation point  Pressure (Pa) Non catalytic heat flux (MW/m2 )  Arc power (kW)  38.0  20  17,500  5.8  2900  0.21  0.55  0.076  0.12  0.046  0  42.5  22  19,300  6.5  3200  0.21  0.55  0.08  0.11  0.05  0  46.0  24  21,000  7.1  3400  0.21  0.53  0.09  0.1  0.07  0  51.0  26  23,000  8.0  3800  0.21  0.52  0.10  0.08  0.08  0  114,000  5  116,000  6  119,000  7  122,000  8  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  1903  Fig. 5. Temperature profiles vs. time during arc-jet testing of the (a) HB5 and (b) HC5-I and HC5-II models corresponding to specific total enthalpies during the test.        for the emissivity at the highest temperature. During the first test on HfC model (HC5-I), the surface temperature achieved by the sample was 2050 C (Fig. 5b). During the second test (HC5-II) the surface temperature reached the value of 2400 C that was maintained for about 4 min. The corresponding computed surface heat flux was of the order of 10 MW/m2 . In both tests a value of about 0.7 was measured for the emissivity at the highest temperatures. The emissivity values of the samples, tested at different conditions, were found to be independent of the test temperature and conditions. These values were similar to those found for other ZrB2-SiC ultra-high temperature ceramics tested in similar conditions.17,22 Both the boride and carbide materials showed an excellent stability during the tests, despite the unavoidable microstructural changes occurring on their surface, as described below.  tion reaction of the two constituent phases. Hafnium diboride oxidises according to23 : HfB2 (s) + (5/2)O2 (g) = HfO2 (s) + B2O3 (l) B2O3 (l) = B2O3 (g)  (2)  (1)           Hafnia is a very stable phase in oxidising atmosphere above 2000 C. It has a melting point of 2900 C and relatively low vapour pressure.24 Boron oxide has a low melting point and high vapour pressure, therefore at T > 1100 C, it starts to evaporate, according to reaction (2). On the other hand, at temperatures > 1000 C, MoSi2 is known to form a stable silica layer according to25 : MoSi2 + (7/2)O2 (g) = MoO3 (g) + 2SiO2  (3)     3.3. Microstructural modifications induced by high enthalpy plasma  3.3.1. HfB2 sample  Due to high thermal loading and the presence of oxidising species, oxidation of the constituent phases occurred. The sample surface was covered by a compact silica-based scale (15-20 \\u242em thick), which embedded HfO2 crystals (Fig. 6a). The composition of the outer glassy layer was investigated by means of EDS analysis. The silica-based scale contained several impurities, including Hf, Al and boron. Bubble formation is due to evolution of gaseous products. The shear forces associated to the hot stream enhanced the bursting of bubbles. The analysis of the cross section (Fig. 6c) revealed that the scale was a multilayered oxide (70 \\u242em thick) well adherent to the bulk, implying that no micro/macrospallation phenomena occurred. Underneath the surface silica oxide (about 10 \\u242em), the scale was mainly constituted by large HfO2 grains and silica with composition close to SiO2 (Fig. 6e). Occasionally Al impurities were detected inside the glassy phase. In the innermost layer formation of molybdenum oxides, silicon oxides and hafnium oxide was observed. Nitrogen impurities were also detected in the composition of the molybdenum-based oxide. The chemistry of the experiment carried out on HfB2 -based material was governed by the oxida The nature of the oxide observed suggests that during the plasma torch tests the silica production according to reaction (3) was fast enough to protect the material. The formation of bubble was produced by the escape of gaseous by-products, such as MoO3 , B2O3 and other highly volatile boron suboxides. However, the presence of B impurities in the outer glassy layer suggests that due to the short exposure the release of boron gaseous products was not complete. The role of dissociated oxygen, which is the primary oxidant agent in a re-entry environment, is still matter of debate. Substituting atomic oxygen to molecular oxygen, thermodynamic calculations (HSC Chemistry for Windows 5, Outokompu Research, OY, Pori, Finland) indicate that reactions (2) and (3) are even more favoured (the absolute value of free Gibbs energy increases by a factor comprised between 1.4 and 1.7). Preliminary results reported in the literature on ZrB2 -SiC composites subjected to re-entry simulations confirmed that oxidation by atomic oxygen proceeds more rapidly than by molecular oxygen.4 Post-tests analyses unfortunately do not help to clarify these aspects. The morphology of the cross section, despite the drastic differences in the experimental conditions, is very similar to that of HfB2-MoSi2 composites oxidised under conventional conditions, such as static air, longer holding times (15 min) and C. Recent studies11 confirm that temperatures of 1400-1650 the response of ZrB2-SiC and HfB2-SiC ceramics to arc heater     \\x0c', '1904  R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  Fig. 6.  (a) HfB2 sample after arc-jet test, (b) surface, (c) cross section, (d) enlarged view of layers I, II and (e) Enlarged view of layer III.  testing appear to be similar to conventional oxidation studies. Finally, it is worth noting that differently from ZrB2 -SiC and HfB2 -SiC composites no depletion layer was observed in the scale, i.e. no active oxidation of MoSi2 occurred for this system.  3.3.2. HfC samples  After the HFC-I tests no significant variation of the sample size and shape was observed, implying that the extent of ablation was very low. The surface turned from dark grey to a whitish colour (Fig. 7a and b) and was constituted by a visibly cracked hafnium oxide scale, with no glassy phase. The specimen section (Fig. 7c) displayed the formation of a multilayered scale, with thickness of about 90 \\u242em, well adherent to the unreacted bulk. Underneath the surface, the outermost portion of the scale was constituted by hafnium oxide and a silica-based glassy phase, which partially filled the porosity (Fig. 7d). The intermediate layer contained a fine porosity and was constituted by hafnia and isolated pockets of molybdenum oxide (Fig. 7e). The innermost layer contained partially oxidised HfC and SiC, and residual Mo-Si species (Fig. 7f). No porosity was found in this region. Similar features were displayed by the model tested at 2400 C (test HfC-II), even if the oxide scale (300 \\u242em thick) was found to be partially detached from the unreacted bulk. Despite the presence of MoSi2 as a SiO2 -forming phase, no evidence of continuous glassy layer was found on the surface of HfC, in contrast with results obtained for HfB2 . This finding cannot be totally imputed to the extreme conditions of the present experiments. The two compositions, HB5 and HC5, were in fact oxidised at the same conditions, in static air (i.e. in absence of        significant ablation) at 1650 C, in a conventional furnace. Even under these milder conditions, HfB2 displayed the formation of a compact layer of silica, while the surface of HfC was mainly covered by an HfO2 scale with few discontinuous pockets of silica. This experiment suggests that the presence of a carbide matrix rather than a boride apparently hindered the formation of a stable silica layer. It can be hypothesized that CO species deriving from oxidation of HfC interacted with MoSi2 causing release of volatile SiO. In the cross section, the layered structure of the oxide (Fig. 7) resemble the results presented in the literature on the oxidation behaviour of monolithic HfC24,26 These studies reported the formation of a layered scale, which comprised a porous outer layer, a dense interlayer and an oxycarbide layer, HfO2−xCy .24,26 The addition of MoSi2 in the composite of the present work resulted in formation of silica which partially filled the inner porosity of the HfO2 scale, a feature which should improve its oxidation resistance. In the innermost layer, HfOxCy species and SiOxCy species were observed, the latter at the interface between MoSi2 and HfC. The formation of SiOxCy species could to be related again to interaction of CO species with MoSi2 at the very low oxygen partial pressure existing under the scale. The oxidation mechanisms of this composite still have to be completely understood and will be object of further investigation. As for the diboride system, the role of atomic oxygen is unclear. In the literature no studies can be found on this topic. The microstructural features of the HfC-based material after testing in the arc-jet facility are indeed very similar to those of samples oxidised in conventional furnaces.  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  1905  Fig. 7.  (a) HfC sample after arc-jet  test, (b) surface, (c) cross section. Enlarged view of (d) layer I (SE imaging), (e) layer II (BSE imaging), (f) layer III (BSE  imaging).  4. Numerical rebuilding for surface catalytic efficiency evaluation  Numerical computations were carried out using the model described in Section 2.3, under different assumptions about the catalytic properties of the specimen surface with reference to catalytic efficiency of atomic nitrogen and oxygen. The simulations refer to the four plasma torch test conditions reported in Table 2. Based on the computed heat flux distributions, a thermal analysis was carried out to evaluate the catalytic efficiency value γ needed to fit the experimental temperature data. The  catalytic efficiency is defined as the ratio of the number of dissociating atoms that recombine at the wall to the total number of the colliding atoms with the wall. For a non catalytic wall this value is 0, while for a fully catalytic wall this is 1. Fig. 8 shows the steady state results computed under the two assumptions of fully catalytic and non catalytic wall and the experimental data obtained with the pyrometer for both material samples. For thermal analysis, input values of specific heat and thermal conductivity are necessary. As a first approximation, the values of monolithic HfC and HfB2 were considered for the composites, as follows: 200 J/(kg C) and 22 W/m C for specific heat and        Fig. 8. Experimental results and numerical solutions corresponding to the different assumption of fully catalytic (FC) and non catalytic (NC)wall for (a) HfB2 and  (b) HfC.  \\x0c', '1906  R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  5. Summary and future work           Two different ultra-high temperature ceramics, HfB2 + 5% MoSi2 and HfC + 5% MoSi2 were produced by pressureless sintering. Machined hemispherical models were exposed to ground simulated atmospheric re-entry conditions using arc-jet testing, with an average specific total enthalpy of the flow around the body of the order of 5-10 MJ/kg and at atmospheric pressure. The HfB2 + 5% MoSi2 model surface reached a peak value of 1950 C for H approaching 8 MJ/kg. SEM-EDS analysis of the cross section after exposure showed the formation of a compact silica oxide (about 15 \\u242em) which sealed the underlying HfO2 scale. The HfC + 5% MoSi2 model surface reached peak values of 2100 C and 2400 C. Cross section analysis showed a layered structure, constituted of an outer layer of porous HfO2 and an inner layer mainly constituted of HfO2 and silica. Numerical calculations, which simulated the chemical nonequilibrium flow around the hemispheric model correlated well with the experimental results assuming a very low catalytic surface behaviour for HfB2 and a catalytic behaviour increasing with temperature for HfC. Although more testing is necessary to improve our understanding of the oxidation mechanisms under extreme conditions, the composites presently tested showed an excellent resistance to high enthalpy hot flows. This stability at temperature around 2000 C opens up new developments in several fields of application, including nuclear applications and industries where extreme conditions are involved. Moreover, the possibility to produce near net shape components through a conventional sintering technique represents a technological advantage in comparison with materials in need of pressure-assisted techniques and expensive post-sintering machining.     References  1. Upadhya, K., Yang, J. M. and Hoffman, W., Materials for ultrahigh temperature structural applications. Am. Ceram. Soc. Bull., 1997, 76(12), 51-  56.  2. Marschall, J., Chamberlain, A., Crunkleton, D. and Rogers, B., Catalytic  atom recombination on ZrB2 /SiC and HfB2 /SiC ultrahigh-temperature 511 ceramic composites. J. Spacecraft Rockets, 2004, 41(4), 576-581.  3. Gasch, M., Ellerby, D., Irby, E., Beckman, S., Gusman, M. and Johnson,  S., Processing, properties and arc-jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics. J. Mater. Sci., 2004, 39, 5925-5937.  4. Bongiorno, A., Forst, C.  J., Kalia, R. K., Li,  J., Marschall,  J., Nakano,  A. et al., A perspective on modelling materials in extreme environments: oxidation of ultrahigh-temperature ceramics. MRS Bull., 2006, 31, 410-  418.  5. Richet, N., Lespade, P., Goursat, P. and Laborde, E., Oxidation resistance  of HfB2 -SiC coatings for protection of carbon fiber based composites. Key Eng. Mater., 2004, 264-268(TTP), 1047-1050.  6.  Janowski, R., Tauche, M., Scheper, M., Monti, R. and Savino, R., Space plane:  a  new way  for  atmospheric  reentry.  In Proceedings  of  the  1st  Inter-national ARA Days, Atmospheric Reentry Systems, Missions and Vehi cles. Session 15-System Design, 2006.  7. Monti, R., De Stefano Fumo, M. and Savino, R., Thermal shielding of a  reentry vehicle by ultra high temperature ceramic materials. J. Thermophys. Heat Transfer, 2006, 20(3), 500-506.  8. Campbell, I. E. and Sherwood, E. M., ed., High-Temperature Materials and  Technology. Wiley, New York, 1967.  Fig. 9. Numerical evaluation of coefficient of catalytic recombination for HfC 5% MoSi2 as a function of the temperature.        thermal conductivity of HfC,27 300 J/(kg C) and 80 W/m C for specific heat and thermal conductivity of HfB2 .28 The data displayed in Fig. 8a and b, at the same plasma torch conditions, i.e. the same free stream conditions, highlight that the heating behaviour of the two materials was different. The experimental results of HfB2 sample matched well the numerical values corresponding to the non catalytic wall condition. This points out that the material herein tested exhibits a non catalytic behaviour at very high temperatures. This behaviour can be explained by the formation of a silica surface layer (Fig. 6) which is known to possess very low catalytic recombination behaviour.29,30 The presence of such a surface layer also justifies the high values of the surface emissivity, according to the literature data. The experimental results for HfC suggested a partially catalytic behaviour. In order to identify a dependence of surface catalytic efficiency with temperature, different values were considered for each test conditions. Fig. 9 shows the results of the computations. At 1800 C HfC exhibited a non catalytic behaviour. increased up to a value of 2 × 10 Increasing the temperature the catalytic efficiency −3 at 2400 C (Fig. 4a and b) which is relatively low with respect to the fully catalytic wall condition (γ = 1) and of the same order of other low catalytic materials such as those of the Space Shuttle tiles.31 It should be pointed out that the present tests have been carried out at atmospheric pressure conditions. Experimental and theoretical works on the catalytic activity of silica-based materials under simulated re-entry conditions32,33 showed that at constant temperature the catalytic atomic recombination coefficients are decreasing functions of the pressure. Therefore the catalytic properties of the material, in respect to the recombination of oxygen atoms, may be larger at lower pressures, as found for instance in arc-jet experiments with ZrB2 /SiC and HfB2 /SiC ceramic materials.2        \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 28 (2008) 1899-1907  1907  9. Clougherty, E. V. and Kaufman, L., Investigation of boride compounds for  high temperature ceramic composites. J. Spacecrafts Rockets, 2006, 43,  very high temperature applications, ManLabs,  Inc., Cambridge, MA, Air  1004-1012.  Force Technical Documentary Report No. RTD-TDR-63-4096, 1963.  10. Levine, S. R. et al., Evaluation of ultra-high temperature ceramics for aeropropulsion use. J. Eur. Ceram. Soc., 2002, 22(14-15), 2757-2767.  11. Fahrenholtz, W. G.  and Hilmas, G.  E., NSF-AFOSR  Joint Work shop  on  Future  Ultra-High  Temperature Materials,  National  Sci ence Foundation Workshop, http://web.umr.edu/uhtm/. 12. Opeka, M., Talmy, I. G., Wuchina, E. J., Zaykoski, J. A. and Causey, S.  Arlington,  January  2004.  VA,  13-14  J., Mechanical, thermal and oxidation properties of refractory hafnium and zirconium. J. Eur. Ceram. Soc., 1999, 19, 2405-2414.  23. Berkowitz-Mattuck, J. B., High-temperature oxidation. III-Zirconium and hafnium diborides. J. Electrochem. Soc., 1966, 113, 908.  24. Wang, C. R. and Yan, J. M., Thermal stability of refractory carbide/boride composites. Mat. Chem. Phys., 2002, 74, 272-281.  25. Lohfeld, S. and Sch ¨utze, M., Oxidation behaviour of particle reinforced  MoSi2 composites at temperatures up to 1700 Mater. Corr., 2005, 56(2), 93-97.  C. Part I: Literature review.        26. Bargernon, C. B., Bendon, R. C., Jette, A. N. and Phillips, T. E., Oxidation of     to 2060  C. J. Am. Ceram.  Hafnium carbide in the temperature range 1400 Soc., 1993, 76, 1040-1046.  13. Wuchina, E., Opeka, M., Causey, S., Spain,  J., Cull, A., Routbort,  J. et  27. Loehman, R., Corral, E., Dumm, H. P., Kotula, P. and Tandon, R., Ultra  al., Designing for utrahigh-temperature applications: the mechanical and thermal properties of HfB2 , HfCx , HfNx and ␣Hf(N)). J. Mat. Sci., 2004, 39, 5939-5949.  14. Silvestroni, L. and Sciti, D., Effects of MoSi2 additions on the properties of Hf- and Zr-B2 composites produced by pressureless sintering. Scripta Mater., 2007, 57, 165-168.  15. Sciti, D. and Silvestroni, L., High-density pressureless sintered HfC-based composites. J. Am. Ceram. Soc., 2006, 89(8), 2668.  16.  Jeng, Y. L. and Lavernia, E. J., Review, Processing of molybdenum disilicide. J. Mat. Sci., 1994, 29, 2557-2571.  17. Monteverde, F. and Savino, R., Stability of ultra-high-temperature ZrB2 -SiC ceramics under simulated atmospheric re-entry conditions. J. Eur. Ceram.  Soc., 2007.  18. Park, C., Nonequilibrium Hypersonic Aerothermodynamics.  John Wiley  &Sons, 1990.  High Temperature Ceramics for Hypersonic Vehicle Applications, SAN DIA REPORT SAND 2006-2925, June 2006, Sandia National Laboratories,  Albuquerque, New Mexico, USA.  28. Krajewski, A., D’ Alessio, L.  and De Maria, G., Physic-chemical  and  thermo-physical properties of cubic binary carbides. Cryst. Res. Technol., 1998, 33, 34.  29. Greaves, J. C. and Linnett, J. W., Recombination of atoms at surfaces. Part 5. Oxygen atoms at oxide surfaces. Trans. Faraday Soc., 1959, 55, 1346.  30. Dickens, P. G. and Sutcliffe, M. B., Recombination of oxygen atoms on oxide  surfaces. Part 1. Activation energies of recombination. Trans. Faraday Soc., 1964, 60, 1272.  31. Gnoffo, P. A. and Inger, G. R., Analytic corrections to CFD heating pre dictions  accounting for  changes  in surface  catalysis, Part 2. AIAA 7th  International Spaceplanes and Hypersonic Systems and Technologies Con ference, Norfolk, VA, November 18-22, 1996, AIAA Paper No. 96-4589.  19. Park, C., Review of chemical-kinetic problems of future NASA missions. I. Earth Entries. J. Thermophys. Heat Transfer, 1993, 7(3), 385-398.  20. Park, C., Howe, J. T., Jaffe, R. L. and Chandler, G. V., Review of chemical 32. Kolesnikov, A. F., Gordeev, A. N., Vasilevskii, S. A. and Verant, J. L., Pre dicting catalytic properties of SiC material  for  the Pre-X vehicle reentry  conditions. In Proceedings of the First European Conference for Aerospace  kinetic problems of future NASA missions. II. Mars entries. J. Thermophys. Heat Transfer, 1994, 8(1), 9-23.  21. Park, C., Jaffe, R. L. and Partridge, H., Chemical-kinetic parameters of hyperbolic earth entry. J. Thermophys. Heat Transfer, 2001, 15(1), 76-90.  22. Scatteia, L., Borrelli, R., Casentino, G., Beche, E.,  Sans,  J. L.  and  Sciences (EUCASS), 2005.  33. Kolesnikov, A. F., Yakushin, M.  I., Pershin,  I. S., Vasilevskii, S. A.,  Chaot, O., Vancrayenest, B. et al., Comparative study of surface catalyt icity under subsonic air  test conditions.  In Proceedings of  the 4th Europ.  Symp. Aerothermodynamics  for Space Applications, 2001, pp. 481-486,  Balat-Pichelin, M., Catalytic and radiative behaviours of ZrB2 -SiC ultra ESA SP-487.  \\x0c']"
},{
  "_id": 19,
  "PDF": "Behavior of HfB2–30 vol_ SiC UHTC obtained by sol–gel approach in the supersonic airflow.pdf",
  "Text": "['Journal of Sol-Gel Science and Technology (2019) 92:386-397  https://doi.org/10.1007/s10971-019-05029-9  O R I G I N A L P A P E R : N A N O S T R U C T U R E D M A T E R I A L S ( P A R T I C L E S , F I B E R S ,  C O L L O I D S , C O M P O S I T E S , E T C . )  Behavior of HfB2-30 vol% SiC UHTC obtained by sol-gel approach in the supersonic airflow  Elizaveta P. Simonenko Vladimir G. Sevastyanov1  1  ● Nikolay P. Simonenko ● Nikolay T. Kuznetsov1  1  ● Andrey N. Gordeev2  ● Anatoly F. Kolesnikov2  ●  Received: 13 December 2018 / Accepted: 20 May 2019 / Published online: 25 June 2019 © Springer Science+Business Media, LLC, part of Springer Nature 2019  Abstract  Using the sol-gel process, a composite powder of HfB2-(SiO2-C) composition was synthesized and then utilized for reactive sintering (hot pressing, 30 MPa; 1800 °C, 15 min) of ultra-high-temperature ceramic HfB2-30 vol% SiC composites. It was established that nanocrystalline silicon carbide (36 ± 2 nm) in a cubic modification is formed during hot pressing; HfO2 and HfC impurities were not found. The relative density of the obtained materials was 89.2 ± 2.3%. The long-term (40-min)  oxidation resistance of the HfB2-30 vol% SiC sample was studied under the exposure to the supersonic dissociated airflow on a high-frequency induction plasmatron (heat flux changed from 232 to 779 W/cm2, pressure in the plasmatron chamber  was 15 hPa),  in a configuration that prevented significant heat discharge from the sample to a water-cooled holder (with 1 : , ;  ) (  0 9 8 7 6 5 4 3 2 1  : , ;  ) (  0 9 8 7 6 5 4 3 2 1  mm overhang). It was shown that, with a stepwise increase in the heat  over the front surface of the sample was relatively uniform. However, at  load in the initial stages, the temperature distribution the heat flux of 598 W/cm2,  local overheated areas  appeared in the central  region, which spread over almost  the entire surface of  the sample within 1-2 min;  the average  temperature was ~2560 °C. Using emission spectroscopy data from the boundary layer above the sample surface, as well as  XRD and SEM of  the sample after exposure,  it was shown that  the sharp increase in temperatures from ~1500-1600 to  2500-2600 °C was associated with a change in the chemical nature of the surface, due to evaporation of borosilicate glass  components and the appearance of a porous highly catalytic HfO2 with low thermal conductivity on the surface. It was noted that these processes under the exposure to the supersonic flow started at lower temperature than under the exposure to  subsonic dissociated airflows.  Graphical Abstract  Using the sol-gel process, a composite powder of HfB2-(SiO2-C) composition was synthesized and then utilized during the reactive sintering (hot pressing, 30 MPa; 1800 °C, 15 min) of ultra-high-temperature HfB2-30 vol% SiC ceramic composites. The long-term (40 min) oxidation resistance of the HfB2-30 vol% SiC sample was studied under the exposure to the supersonic dissociated air flow on a high-frequency induction plasmatron (heat flux changed from 232 to 779 W/cm2,  pressure in the plasmatron chamber was 15 hPa), in a configuration that prevented significant heat discharge from the sample  to a water cooled holder  (with 1 mm overhang).  * Elizaveta P. Simonenko ep_simonenko@mail.ru  1  Kurnakov Institute of General and Inorganic Chemistry of the Russian Academy of Sciences, Leninsky pr., 31, Moscow 119991, Russia  2  Ishlinskii Institute of Problems of Mechanics of the Russian Academy of Sciences, pr. Vernadskogo, 101-1, Moscow 119526, Russia  \\x0c', 'Highlights  ●  HfB2-(SiO2-C) composite powder was synthesized using the sol-gel process. HfB2-30 vol% SiC UHTC was produced by reactive sintering of HfB2-(SiO2-C) powder. Oxidation resistance of UHTC under the exposure to the supersonic airflow at 2500 °C was studied.  ●  ●  ●  Features of material oxidation were identified,  in particular utilizing emission spectroscopy of  the boundary layer.  ●  The composition and microstructure of  the oxidized material  layer were studied, both its surface and thin section.  Keywords  Sol-gel processes ● Nanomaterial  ● Ceramics ● UHTC ● Supersonic airflow ●  Induction plasmatron  1 Introduction  Ultra-high-temperature ceramic materials (UHTC) based on  zirconium or hafnium diborides modified with silicon car bide are attracting increasing attention [1-15] as materials  capable of withstanding the impact of high-enthalpy airflow  without catastrophic destruction, despite the fact that surthan 2500-2700 °С.  face  temperature  can become higher  Owing to low self-diffusion and strong covalent bonds  in  the crystal  lattices of ZrB2 and HfB2, high temperatures are required during hot pressing, spark plasma, or pressureless  sintering  to  obtain  high-density UHTCs  utilizing  such  compounds. This  results  in a  significant growth of MB2 (where M is Zr, Hf) and SiC grains, and consequently, in  deterioration  of mechanical  properties. Thus,  there  is  an  urgent need to maximally reduce production temperatures  of ceramic materials based on ZrB2(HfB2)-SiC. The use of nanoscale or submicron SiC in the production  of ultra-high-temperature ceramic materials allows not only  for  improvements  in the process of material densification  [16,  17],  but  also  for  significant  improvements  in  their  mechanical properties [17-24]. So, for example,  it  is stated  in ref. [22] that when using SiC powder with a particle size  of  30 nm,  an  increased  resistance  to  thermal  shock  is  observed in the  resulting ZrB2-20 vol% SiC material. experimentally shown in ref. [17] that  In  addition,  it  is  increasing the proportion of nanosized SiC powder in the [(1-x)SiC+xnanoSiC]  composition  of  ZrB2-20 vol%  materials leads to an almost twofold increase in thermal conductivity in the temperature range of 20-1100 °С, and  this is also noted in ref. [23]. Especially important data are  presented in the work of Kovacova et al.  [24], where it  is  shown that  for ZrB2-20 wt.% SiC ceramics, during hot pressing from SiC powder with an average size  the transition  of 44 μm to nanoand submicron size leads to a significant  increase  in  oxidation  resistance when  calcined  in  air  at  temperatures of 1500 and 1700 °C.  Sol-gel  technology is extremely useful  for synthesis of  refractory carbides in the highly dispersed state,  including  silicon carbide [25-35], since the production of intermediate MOx-C compounds (M = Si, Ta, Ti, Zr, Hf, etc.), the components of which are maximally dispersed and  evenly distributed among themselves, makes it possible to  reduce the temperature of carbothermic synthesis.  In the field of ultra-high-temperature ceramic materials  production,  sol-gel  technology is being successfully used  only in the synthesis of highly dispersed ZrB2(HfB2)-SiC composite powders [36-44], which can then be used in the  processes of hot pressing or slip casting of ceramics, as well  as  for  applying  protective  antioxidizing  coatings  on  the  surface of graphite, Cf/C, or Cf/SiC composites. Previously [45], we proposed combining stages of  the  carbothermic SiC synthesis and the consolidation of UHTC,  since the temperatures of  these processes are similar, and  thus to proceed to the technology of reactive sintering of the  intermediate product — HfB2-(SiO2-C) composite powder  Journal of Sol-Gel Science and Technology (2019) 92:386-397  387  \\x0c', 'produced by the sol-gel process — at  relatively low tem peratures.  In  addition,  in  ref.  [45], we  have  noticed  the  increased oxidation resistance of UHTCs produced by this  method at  temperatures of 1700-1900 °C in airflow under  the conditions of a DSC/DTA/TGA experiment.  The study of  the oxidation resistance of  the ultra-high temperature  ceramic materials  obtained  by  the  proposed  method is extremely important, because it makes it possible  to  evaluate  their  applicability  in  the  intended  operating  conditions. Despite  the  large  number  of  useful works  [7,  9,  10,  13,  24,  45-54], which  describe  the  thermal  behavior  of UHTCs  of ZrB2(HfB2)-SiC composition aerodynamic heating by high-speed airflow is  in  static  air,  more  correctly modeled  using  oxyacetylene  torches,  in  electrical  arc units, using induction plasmatrons, or  in a  chamber of a hypersonic ramjet engine [8, 14, 15, 55-65].  For materials of HfB2-30 vol% SiC composition, produced at the minimum temperature of 1700 °C by hot  pressing of HfB2-(SiO2-C) density of 85.8%, behavior under the exposure to sub[64]  composite powders, having a  and supersonic [65] dissociated airflow has been studied on  a high-frequency induction plasmatron. The special  feature  of  the experimental setup was that  the samples were flush  mounted in a water-cooled copper holder, which under  the  exposure  to supersonic flow resulted in the  surface  tem perature not exceeding 1600 °C due to the heat  transfer  to  the holder.  The purpose of this work is to study the behavior of the  HfB2-30 vol% SiC ultra-high-temperature ceramic composite produced by hot pressing of a highly dispersed HfB2- (SiO2-C) composite powder at a higher temperature of 1800 °C, under  the exposure to the supersonic dissociated  airflow, using a high-frequency induction plasmatron.  2 Experimental section  Reagents  used:  tetraethoxysilane  (TEOS)  Si(OC2H5)4 (>99.99%, EKOS-1 JSC), LBS-1 bakelite varnish (Karbolit OJSC), formic acid СН2О2 (>99%, Spektr-Chem LLC), and hafnium diboride (>98%, particle size 2-3 microns, aggre gate size ~20-60 microns, Tugoplavkie Materialy LLC).  HfB2-(SiO2-C) using sol-gel technology,  composite  powder  was  synthesized  according  to  a  previously  described  technique  [37,  45]  (Fig.  1). To  do  this,  tetra ethoxysilane and the hydrolysis catalyst, formic acid, in the n(Si):n(C) = 1:3.05 n(Si):n(CH2O2) = 1:4, phenol-formaldehyde resin  ratios  of  and  were  introduced  dropwise  to  solution (LBS-1 bakelite varnish, carbon source after pyr olysis),  the solution being stirred at  the same time. After  heating the  solution to 40-50 °C, TEOS hydrolysis with  water was initiated; the amount of water corresponded to the ratio of n(Si):n(H2O) = 1:5. HfB2 powder was dispersed in  the resulting colloidal solution. After gelation and stepwise  drying at  the temperature of 50-80 °C, the xerogel was heat  treated  at  the  temperature  of 400 °C under ~1-5•10-6 atm).  a  dynamic  vacuum (residual  pressure  In  this  case,  HfB2-(SiO2-C) the components were distributed as evenly as possible, and  composite powder was  formed,  in which  the SiO2-C compound had an increased reactivity [37]. The production of HfB2-30 vol% SiC ultra-hightemperature ceramic materials was carried out using a hot  press of Thermal Technology Inc.  (HP20-3560-20)  [45].  HfB2-(SiO2-C) composite powder obtained using sol-gel technology was loaded into a graphite mold, compacted,  and evacuated, and then the chamber was filled with argon.  After  this, heating began at  the  rate of 10 °C/min to the  temperature of 1800 °C, and the pressure was adjusted to  the  target value  (30 MPa). The holding time  at  the max imum temperature was 15 min. A small amount of boron  nitride was used as a mold lubricant.  The X-ray diffraction patterns of  the synthesized com posite powders were recorded on a D8 Advance (Bruker) Xin the range of 2Θ 34‒37° with a resolution of 0.02°, the signal being accumulated for 2 s at the point, and in the range of 2Θ 5-80° with a resolution of 0.02°, the signal being accumulated for 0.3 s at the point.  ray diffractometer  The  infrared  reflectance  spectra  of  ceramic materials  were  recorded using  an InfraLUM FT-08 FTIR spectro meter  (PIKE EasiDiff diffuse reflectance accessory).  Scanning electron microscopy (SEM) data were obtained  with a triple-beam workstation NVision 40 (Carl Zeiss);  the  elemental  composition  of microdomains was  determined  with an EDX system (Oxford Instruments).  Fig. 1 Synthesis scheme for HfB2-(SiO2-C) composite powder  388  Journal of Sol-Gel Science and Technology (2019) 92:386-397  \\x0c', 'Experiments  on  the  effects  of  supersonic  dissociated  airflow on the sample surface were performed on a 100-kW  high-frequency  induction  plasmatron  IPG-4  [6, 14, 15, 64, 65]  at  IPMech RAS with an anode  feed  power of the plasmatron (N) from 20 to 70 kW using a sonic  nozzle with an exit diameter of 30 mm. The airflow rate was  3.6 g/s, and the pressure in the chamber was 15 ± 1 hPa. The  temperature of  the  sample  surface was measured using a  Mikron M-770S pyrometer  in the pyrometer  spectral-ratio  mode  (temperature  range of 1000-3000 °C, measurement  spot  size~5 mm). To  record  the  temperature  distribution  over  the sample surface, a Tandem VS-415U thermal  ima ger was used: the measurements were carried out at the set value of the spectral emissivity of ε = 0.6 at a wavelength of 0.9 μm. ε was  The  value  of  selected  in  comparative  experiments on the simultaneous determination of  the sur face temperature of similar  samples, using a Tandem VS 415U thermal  imager and a Mikron M-770S spectral-ratio  pyrometer in such a way as to ensure equal readings on both  devices.  To record the emission spectra of the boundary layer, we  used a HR-4000 high-resolution compact diffraction spec trometer  (Ocean Optics, USA) with a linear CCD detector  (3648  pixels)  and  fiber-optic  radiation  input.  In  this  experiment,  the  device  recorded  a  spectral  range  of  200-650 nm. The  optical  axis was  parallel  to  the  front  surface of the sample in the boundary layer. The stagnation pressure Pst was measured using a Pitot tube with a hemispherical nose (R = 15 mm) and a  receiving-hole diameter of 14 mm.  3 Results and discussion  3.1 Study of the HfB2-SiC samples obtained by sol-gel technique  The density 7.89 g/cm3  of  the  obtained  HfB2-SiC value calculated  samples  was  (89.2 ± 2.3% of  the  by  the  additive method, where the density of HfB2 was taken to be 11.2 g/cm3 [66], and the density of SiC was 3.2 g/cm3 [67]).  X-ray  diffraction  analysis  of  the  obtained  samples  (Fig. 2) suggested that  the synthesis of silicon carbide took  place in full—against the background of intense reflexes of there were reflexes of the β-SiC phase the HfB2 phase [68], [69], but neither reflexes of the crystalline forms of SiO2 nor a diffuse halo associated with the content of the amorphous  SiO2 phase were found. nium oxide (as a product of HfB2 oxidation) or hafnium carbide was not observed, which, as noted in the literature,  In addition,  the formation of haf can  be  synthesized  by  hot  pressing  if  the  initial HfB2 powder has a high content of impurity oxygen. The average  crystallite size of synthesized SiC estimated by the Scherrer  equation was 36 ± 2 nm.  The diffuse reflectance IR spectroscopy data confirmed a  complete conversion of SiO2 bers in the 970-1200 cm-1 range,  into SiC:  for  the wave num there was no broadened  intense absorption band corresponding to stretching vibra tions of  the Si-O groups that  is characteristic of  the initial  powders of  the HfB2-(SiO2-C) composition; however, absorption band in the range of 800-950 cm-1 that characteristic of ν(Si-C) did appear.  the  is  3.2 Behavior of the prepared HfB2-SiC samples in supersonic dissociated air flow  To study the behavior of UHTC under heating conditions,  under  the  exposure  to the  supersonic dissociated airflow  using  an  induction  plasmatron  IPG-4,  the  sample was  placed in a copper model  (Fig. 3), consisting of a mandrel  (into which the sample was directly inserted) and a water cooled holder into which the mandrel moves along a sliding  fit. Thermal contact on the end surface was guaranteed by  springing the tension pin; contact surfaces were lubricated  with thermal grease,  for  improved heat  transfer  from the  mandrel  to  the water-cooled  holder.  The  sample was  installed in the mandrel using three whisker wool filaments  based on the SiC fiber-like crystals with a 1-mm protrusion  (Fig. 3b),  to reduce heat  transfer  from the  sample  to the  model compared with the fixing layout described in ref. [65]  (Fig. 3a). A conical sonic nozzle with an exit diameter of  30 mm was used, from which the distance to the sample was  25 mm.  Fig.  2 X-ray diffraction pattern of HfB2-30 vol% SiC samples obtained by hot pressing of HfB2-(SiO2-C) composite powder at a temperature of 1800 °C; the insets show in more detail the 2Th intervals, in which the most intense reflexes of the phases of SiC [69] (a) as well as of HfO2 [71], SiO2 [72], and HfC [73] (b) are observed  Journal of Sol-Gel Science and Technology (2019) 92:386-397  389  \\x0c', 'The sample was introduced into a dissociated airflow at an anode feed power of the plasmatron N = 20 kW and with the pressure in the chamber Pch = 15 ± 1 hPa. Further, power was increased to 70 kW with an increment of 10 kW. The aging duration at N = 20-60 kW was 3 min, and the total exposure time—40 min. The value of heat flux (q) to the high catalytic copper surface increased from 232 (N = to 779 W/cm2 (N = 70 kW)  the  20 kW)  (Table 1).  The change in the average surface temperature measured  by the spectral-ratio pyrometer is presented in Fig. 4 and in  Table 1 as a function of  the impact parameters. (N = 20-50 kW),  In the first  stage of  the experiment  the average tem perature increased with the heat flux, with a slight tendency to decrease (~10 °С) during aging. However, at N = 60 kW (q = 598 W/cm2),  the  surface  heated  by  almost  100 °C  within 3 min, as a result of which the temperature exceeded  1800 °C. A further  increase in N led to a sharp increase in  the average surface temperature up to ~2560 °C, which then  slightly  decreased  (to  2540 °C)  during  prolonged  aging  (25 min).  The analysis of  thermal  images (Fig. 5)  recorded by the  thermal  imager allows for a more detailed understanding of  the processes occurring on the  surface of  the HfB2-SiC sample when interacting with a high-enthalpy airflow. As  can be seen, once the sample was introduced into the dis sociated airflow,  the temperature distributed over  the sur face fairly evenly. This is more evident from Fig. 6, where  the temperature distribution of the surface along the sample  diameter marked by the line in Fig. 5 is shown.  Increasing  the anode feed power of  the plasmatron to 50 kW led to a  stepwise increase in the entire surface temperature, but at N = 60 kW,  small  areas with  a  diameter  of  ~0.5-1 mm  began to form on the surface near  the central  region of  the  sample  by  the  15th minute;  their  temperature was  sig nificantly higher  than the average surface temperature.  With an additional  increase in power up to 70 kW (in  accordance with the heating mode),  the area of  the over heated  sites  increased  drastically  within  1-2 min  and  occupied almost the entire surface of the sample; perature increased to ~2560 °С (Figs. 5 and 6).  the tem It  should  be  noted  that  the  highest  temperature was  observed  in  the  center  of  the  sample,  as was  noted  for  applying supersonic flow to HfB2-SiC samples in ref. (though in that study it was in a different temperature  [65]  range),  in contrast  to the situation typical of the exposure to  subsonic  dissociated  airflow [6,  14,  15,  64], when  the  increase in the surface temperature spread from the edges of  the sample. For  the experiment presented, despite the fact  Fig. 3 Sketch of the model into which the samples were installed for testing: sample fixing diagram in ref. [65] a and in this work b  Table 1 The change in the average temperature of the sample surface (spectral-ratio pyrometer, Тpyr) as a function of exposure time and process parameters: anode feed power (N) and pressure in the plasmatron chamber (Pch), as well as the corresponding heat flux (q) and stagnation pressure (Pst)  Time, min  N, kW  Pch, hPa  q, W/cm2  Pst, hPa  Тpyr, °С  0→3  20  15.0  232  50.2  1249→1230  3→6  30  15.3  363  55.8  1398→1393  6→9  40  15.1  484  59.6  1506→1513  9→12  50  15.1  598  62.5  1644→1632  12→15  60  14.9  691  65.8  1731→1824  15  70  14.9  779  68.7  1951  17  70  14.5  779  68.7  2569  20  70  14.3  779  68.7  2561  25  70  14.3  779  68.7  2566  30  70  14.7  779  68.7  2562  35  70  14.7  779  68.7  2553  40  70  14.5  779  68.7  2543  Fig. 4 Change in the average temperature (according to the pyrometer) of the surface of HfB2-30 vol% SiC sample, depending on anode feed power N, pressure in a plasmatron chamber Pch, and holding duration  390  Journal of Sol-Gel Science and Technology (2019) 92:386-397  \\x0c', 'Journal of Sol-Gel Science and Technology (2019) 92:386-397  391  Fig. 5 Thermal images of the surface of HfB2-30 vol% SiC sample at different points of exposure to the supersonic dissociated airflow; highlighted images show the appearance, and significant increase in, the number of overheated local areas  the  Fig. 6 Temperature distribution along the diameter of HfB2-30 vol% SiC sample at different points of exposure  that  the average surface temperature exceeded from 2500 to  2600 °C,  there remained a small area at  the periphery of the  sample temperature, of which it did not exceed 2200 °C for  40 min (Fig. 7).  The described sharp increase in the surface temperature, as  is known [14, 15, 70], may be associated with a change in the  surface  chemical  composition,  due  to  oxidation  and  eva poration. At  relatively low temperatures  (<1700-1850 °C),  the surface is composed predominantly of borosilicate glass  layer,  and at  temperatures of >1900-2000 °C,  an intense  evaporation not only of boron oxide, but  also of  silicon  oxide takes place, which leads to refractory, highly cataly tic,  and possessing low thermal  conductivity hafnium (or  zirconium) oxide emerging to the surface,  i.e., giving rise  Fig. 7 Change in temperatures in different areas of the HfB2-30 vol% SiC sample in the course of the experiment  surface of  \\x0c', 'essentially to the formation of a thermal barrier  layer  that  prevents the transfer of heat from the overheated sites. The  emission spectroscopy of  the gas phase  in the boundary  layer  above  the  sample  surface  allows  one  to  study  experimentally the change in the intensity of the borosilicate  glass component evaporation (Figs. 8 and 9).  As can be seen from these figures, heating the surface to  1200-1400 °C did not lead to intensive evaporation of either boron oxide or silicon oxide (one and two stages, q = 232 and 363 W/cm2,  respectively); however,  low-intensity  B and Si  lines began to appear in the emission spectrum at  the  average  temperature  of  1500 °C (8b  spectrum). The subsequent increase in the heat exposure to q = 484 and 598 W/cm2 caused a sharp increase in the intensity of  the  boron  and  silicon  lines, which  gradually  decreased  sig nificantly with aging, which may indicate the establishment  of a new equilibrium between the chemical and physico chemical oxidation and evaporation processes on the surtemperatures ≤1500 °C. However, at N = 60 kW (q = 691 W/cm2),  face at  the anode feed  power  of  points  with  increased temperature started appearing on the 14th minute  (probably due  to greater  evaporation of  the  least volatile  glass  component—SiO2), which maximum on the curve 9b. With an increase to the highest level of heat flux q = 779 W/cm2,  corresponds  to  a  local  the  intensity of vapor ization processes was so great that the spectrometer went off  scale. At  the same time,  the surface area of the sample with  a  temperature  of  >2000 °C rapidly  increased  (probably  HfO2 when the size of the high-temperature surface region almost  appeared on the  surface)  (Fig. 5),  and after 2 min,  stabilized,  the  intensity  of  the  boron  and  silicon  lines  decreased to measurable values.  Aging the sample at  the maximum value of heat flux for  ~10 min led to a gradual decrease  in the  intensity of  the  silicon line. A similar  trend was observed for  the intensity  of  the  boron  line,  but was  extended  over more  time  (Fig. 9c). For the last 5 min of exposure with the surface temperature of ~2550 °С,  the intensities of  the boron and  silicon lines were close to those for the initial stages of exposure, when the surface temperature was <1500 °С.  In general, spectral data indicated that, when exposed to  the  supersonic  dissociated  airflow and  at  the  stagnation  pressure of ~60-65 hPa, which was less than for the modes  used to study the effects of the exposure to subsonic airflow at N = 60 kW in this work vs. 100-200 hPa  (65.8 hPa  in  refs.  [14, 15, 64]),  the evaporation of  the borosilicate glass  layer from the sample surface occurred at a noticeably lower  temperature.  The  total weight  loss  of  the  sample  after  40 min  of  exposure was  2.6%. The  increase  in  the  thickness  of  a  cylindrical  sample  of HfB2-30 vol% SiC after  oxidation  was 0.4 mm (11%).  Fig.  8 Emission spectra (wavelength intervals of 246-254 and 286-292 nm) of the boundary layer above the sample surface of HfB2-30 vol% SiC composition depending on the time of the test (according to refs. [74, 75], the sensitive lines of boron are 249.678 and 249.773 nm; of silicon they are 250.690, 251.432, 251.611, 251.921, 252.412, 252.851, and 288.158 nm)  Fig. 9 The average surface temperature of HfB2-30 vol% SiC sample а and the relative intensity of the silicon lines at 251.7 nm b and boron lines at 249.8 nm c, depending on holding time; the “a-f” designation corresponds to the spectra at different points of the experiment, shown in Fig. 8  392  Journal of Sol-Gel Science and Technology (2019) 92:386-397  \\x0c', '3.3 Study of the HfB2-SiC sample after the exposure to the supersonic dissociated airflow  The XRD of the surface after exposure showed that both for  the region with an average temperature of ~2550 °C and for  a relatively “low-temperature” area on the sample edge,  the  temperature of which did not  exceed 2170 °C during the  test,  the phase composition was the HfO2 monoclinic phase (Fig. 10). The microstructure (according to the SEM data)  also differed only slightly (Figs. 11 and 12): on the surface,  there was  a  porous  ceramic  layer, mainly  consisting  of  hafnium dioxide. At  the  same  time,  the  formation of  the  layer with a larger pore size of up to 10 μm is typical for the ratio of n(Hf):n(Si) =  high-temperature areas (Fig. 11);  215:1 (EDX). A relatively “low-temperature” area is char acterized by a greater roughness, and on the surface of HfO2  particles,  there were  traces  of  Si-containing  formations  (Fig. 12c, d), probably remaining after the evaporation of ratio of n(Hf):n(Si) = 9:1  the borosilicate glass melt;  the  (EDX).  The  study  of  the  thin  section microstructure  and  the  mapping of the distribution of the elements (Si, O, and Hf)  indicated (Fig. 13) that, as a result of a 40-min exposure to a  supersonic dissociated airflow, a multilayer oxidized region  was  formed,  as  noted  earlier  in  cases when  the  surface  temperature  exceeded 2000 °C.  In particular,  there was  a  HfO2-based layer on the surface (points 1-3), in the volume of which a relatively thin layer of borosilicate glass was  localized, preventing oxygen diffusion. In the deeper layers  (points 4-6),  there was a layer of material depleted of sili con  carbide,  due  to  its  active  oxidation with  a  reduced  content  of  oxygen  diffused  through  the  glass  layer;  its  thickness was ~600-800 microns. The total  thickness of the  oxidized area of  the sample was ~1-1.2 microns.  4 Conclusions  In  this work,  using  reactive  sintering  (hot  pressing  at  a  relatively low temperature of 1800 °C, 30 MPa, 15 min of  aging) of  the HfB2-(SiO2-C) composite powder obtained sol-gel process, ultra-high-temperature ceramic  by  the  composite materials  of HfB2-30 vol% SiC composition It has been shown that the indicated  have been produced.  parameters of hot pressing are sufficient  for complete con version of silicon oxide into carbide. According to the XRD  data, there were no HfO2 (HfB2 oxidation product) and HfC (product of HfO2 interaction with carbon obtained by pyrolysis of phenol-formaldehyde  resin)  impurities.  The  Fig. 10 X-ray patterns of the surface of HfB2-30 vol% SiC sample after exposure to the supersonic dissociated airflow sites, the temperature of which was ~2500-2550 °C (1) and did not exceed 2170 °C (2)  Fig. 11 The microstructure of the surface of HfB2-30 vol% SiC sample after exposure to the supersonic dissociated airflow at a temperature of ~2500-2550 °C: a, b, c—according to the data of the secondary electron detector, d—in contrast mode by the effective atomic number, SEM  Journal of Sol-Gel Science and Technology (2019) 92:386-397  393  \\x0c', '394  Journal of Sol-Gel Science and Technology (2019) 92:386-397  Fig. 12 The microstructure of the surface of HfB2-30 vol% SiC sample after exposure to the supersonic dissociated airflow at a temperature of ~2000-2170 °С: a, b, c—according to the data of the secondary electron detector, d—in contrast mode by the effective atomic number, SEM  Fig. 13 The microstructure of the thin section of HfB2-30 vol % SiC sample after exposure to the supersonic dissociated airflow (the sample is fixed in epoxy resin, SEM), and the distribution mapping of Si, O, and Hf elements; the ratio of n (Hf):n(Si) estimated by EDX is highlighted in pink  synthesized silicon carbide was nanocrystalline (the average  carbide is equally spaced in the volume of the material, and  crystallite size calculated by Scherrer  technique was 36 ±  dispersion in a liquid medium leads to the destruction of the  2 nm). The density of the obtained samples was 89.2 ± 2.3%  of  the theoretical one.  HfB2 particle aggregates without For the first time ever, the effect of exposure to super the grinding stage.  That means, due to sol-gel  technology, which allows to  sonic  dissociated  airflow on  a  sample  obtained  by  the  obtain the most  reactive SiO2-C system on the surface of the HfB2 particles, allows us to obtain the more energyefficient UHTC composites of HfB2-30 vol% SiC in lesser amount of stages in relatively soft conditions (hot pressing  described technique has been studied in a configuration that  made it maximally difficult  to transfer  the incoming heat  from the sample to the copper water-cooled holder, namely,  the  configuration with  a  1-mm overhang  relative  to  the  at 1800 °C, 30 MPa for 15 min), under which SiC is nano mandrel  face  (Fig.  3b).  It  has  been  shown  that  such  a  crystalline. In this case,  the resulting nanocrystalline silicon  configuration led to the heating of  the sample surface to a  \\x0c', 'temperature of ~2550 °C, which was not observed when the  sample of  the same composition (somewhat more porous)  was placed flush [65].  It  has  been  shown  experimentally  (and  taking  into  account  the data of emission spectroscopy of  the boundary  layer above the sample of HfB2-30 vol% SiC) the influence of the supersonic airflow at the  that, under  stagnation  pressure of 50.2-68.7 hPa, a sharp increase in the surface temperature to 2500-2560 °С associated with the evapora tion of  silicon and boron oxides  from the surface and the  appearance of highly catalytic and low thermal-conductive  porous HfO2 on it occurred at was observed under the influence (Pch = 100-200 hPa) appearance of boron lines before the emergence of silicon  the lower  temperature than  of  the  subsonic flow  [14, 15, 64]. At  the  same  time, no  lines was  observed;  in  the  experiment  performed,  the  appearance and growth in intensities of  the boron and sili con lines occurred simultaneously. This is probably due to  the high chemical activity of nanocrystalline silicon carbide  obtained by the sol-gel method (crystallite size is ~36 nm):  the  beginning  of  its  oxidation  occurs  at  the  lower  tem peratures close to those for HfB2, which gives an opportunity to form a protective layer of the borosilicate glass  earlier  and  prevent  intensive  “low-temperature”  (<1200-1300 °C)  oxidation  of UHTC  accompanied  by  distillation of  fugitive boron oxide.  It has been noted that, as a result of prolonged (40-min)  exposure supersonic flow of  the dissociated air, which is  characterized  by  complete  dissociation  of  the  oxygen  molecules  and  partial  dissociation  of  the  nitrogen mole cules,  there was no catastrophic destruction of  the sample,  and no cracking or detachment of  the oxidized area;  the  total mass loss due to oxidation and evaporation from the  surface was 2.6%.  In general,  the experiment showed promise of  the devel oped method for producing ultra-high-temperature  ceramic  composites of  the HfB2-30 vol% SiC composition proposed for use at elevated temperatures in oxygen-containing media.  Acknowledgements The study has been funded by the Russian Science Foundation (17-73-20181, for obtaining HfB2-SiC ultra-hightemperature ceramic materials with nanocrystalline silicon carbide using the sol-gel process and studying the mechanism of their oxidation), and by the Russian Foundation for Basic Research (No. 1701-00054-a, for studying the heat transfer of high-enthalpy gas jets with the surface of ceramic samples).  Compliance with ethical standards  Conflict of  interest The authors declare that  they have no conflict of  interest.  Publisher’s  note:  Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.  References  1. Simonenko EP, Sevast’yanov DV, Simonenko NP et al. (2013) Promising ultra-high-temperature ceramic materials for aerospace applications. Russ J Inorg Chem 58:1669-1693. https://doi.org/ 10.1134/S0036023613140039 2. Zoli L, Vinci A, Galizia P et al. 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  "PDF": "Behavior of Some Refractory Hafnium and Tantalum Compounds in Plasma Flows.pdf",
  "Text": "['ISSN 0020-1685, Inorganic Materials, 2019, Vol. 55, No. 3, pp. 231-236. © Pleiades Publishing, Ltd., 2019. Russian Text © N.I. Baklanova, V.V. Lozanov, A.A. Kul’kov, E.A. Antipov, A.T. Titov, 2019, published in Neorganicheskie Materialy, 2019, Vol. 55, No. 3, pp. 257-263.  Behavior of Some Refractory Hafnium and Tantalum Compounds in Plasma Flows  N. I. Baklanovaa, V. V. Lozanova, *, A . A . Kul’kovb, E. A . Antipovb, and A . T. Titovc  aInstitute of Solid State Chemistry and Mechanochemistry, Siberian Branch, Russian Academy of Sciences, ul. Kutateladze 18, Novosibirsk, 630128 Russia bCentral Research Institute for Special Machinery Joint Stock Company, Zavodskaya ul., Khotkovo, Moscow oblast, 141371 Russia cSobolev Institute of Geology and Mineralogy, Siberian Branch, Russian Academy of Sciences, pr. Akademika Koptyuga 3, Novosibirsk, 630090 Russia *e-mail: lozanov.25@gmail.com  Received May 24, 2018; revised October 4, 2018; accepted October 15, 2018  Abstract—By reacting tantalum or hafnium carbide with iridium in the presence of a small amount of silicon, we have prepared refractory hafniumand tantalum-containing materials consisting of a mixture of phases: the intermetallic compound MIr3, recrystallized tantalum or hafnium carbide, and iridium silicide. We have studied the behavior of the materials during an exposure to a high-speed plasma f low at a sample surface temperature of 2000°C and demonstrated that, owing to their special microstructure, the absence of pores, and the low oxidation rate of the iridium-containing components, they exhibit a good ablation resistance and that the hafnium system withstands a longer exposure.  Keywords: tantalum carbide, hafnium carbide, iridium, intermetallic phase, oxidation resistance  DOI: 10.1134/S00201685190300 4X  INTRODUCTION  Existing approaches for protecting structures from degradation  in oxidizing media at temperatures of 2000°C and above rely, f irst, on the “self-healing” ability of systems such as Hf(Zr)B2-SiC [1-3]; second, on the use of a so-called high-entropy ceramic: an equimolar mixture of f ive or more refractory carbides or borides capable of preventing oxygen indiffusion into the bulk of the material [4]; and, f inally, on the use of materials having a slow recession rate in oxygen,  including multilayer  iridium coatings  [5]. There is a strong need for novel materials, because existing ones are incapable of ensuring stable operation of structures at temperatures above 2000°C. There  is certain  interest  in systems combining group IV-VI  transition metals; noble metals,  for example, iridium; and light elements, such as boron, carbon, or silicon [6]. Ternary phase diagrams of such systems are not always known. At the same time, it follows from analysis of the phase diagrams of M-Ir binary systems that compounds of transition metals and iridium, MIrn, have melting points above 2000°C [7]. Moreover, they have a high elastic modulus and high thermal conductivity [8].  The most widespread method for preparing compounds of transition metals with iridium is by reacting the metals in an electric arc furnace under an inert gas atmosphere [8, 9]. Another approach is to react transition metal carbides with iridium. As shown earlier [6, 10, 11], reactions of the refractory carbides HfC, TaC, and ZrC with iridium become noticeable even at 1200°C. They yield only MIr3 phases, the richest in iridium, even if there is an excess of the starting carbides relative to iridium [6, 11]. Note that no iridides with other compositions are formed, which can be accounted for by the low reactivity of the transition metal on the Ir/MC interface [12]. It is worth pointing out that it is iridium-rich transition metal intermetallic compounds which have melting points near 2500°C [7]. The presence of the transition metals in the reaction products can be regarded as an advantageous factor because they have very high melting points and their oxidation products (metal oxides) also have high melting points and low vapor pressure [13].  The subjects of this study were two systems, Ta- C-Ir and Hf-C-Ir, which are of interest for ultrahigh-temperature materials research. The purpose of this work was to study the behavior of refractory mate 231  \\x0c', '232  BAKLANOVA et al.  rials based on these two systems during a prolonged (up to 1000 s) exposure to a high-speed plasma f low at a temperature of 2000°C on the sample surface.  EXPERIMENTAL  Green compacts were prepared using TaC (Russian Federation Purity Standard TU 6-09-03-443-77), HfC (Russian Federation Purity Standard TU 6-0903-361-78), and iridium (Russian Federation State Standard GOST 12338-81, 99.96+% purity) powders. The starting powders were mixed in the ratio MC : Ir = 2 : 3. A small amount of metallic silicon (Russian Federation Purity Standard TU 48-4-319-86) was added to the starting mixture in order to obtain an intermediate liquid phase because, according to phase diagram data, the Ir-Si system has a few eutectics with relatively low melting points [9]. The powder mixture was pressed into disks, which were then f ired in a vacuum furnace at 1600°C for 4 h.  The resultant samples were characterized by X-ray diffraction on a D8 Advance powder diffractometer (Bruker, Germany) at room temperature (CuKα  characteristic radiation, angular range 10° < 2 θ  < 90°). Qualitative analysis of the X-ray diffraction data was carried out using Search-Match software (Oxford Cryosystems) and ICDD PDF-2 Powder Diffraction File data (release 2008).  The surface morphology of the compacts and their cross sections was examined by scanning electron microscopy (SEM) on a TM-1000 (Hitachi, Japan) and MIRA 3 LMU (TESCAN, Czech Republic) equipped with  INCA Energy 450 XMax 80 and Swifted-TM energy dispersive X-ray  spectrometer (EDS) systems, respectively. To determine the elemental composition of the compacts, they were cut, potted in epoxy, and polished on a Struers Tegramin30 polishing machine with 1-μm diamond paste (MonoSyn Duo, Synercon, Germany). Next, a conductive layer was produced on the samples.  The oxidation resistance of the samples was studied using an EDG-200M plasma source. The f lat sample surface was perpendicular to the plasma jet. The working current was varied from 120 to 400 A and the applied voltage was 120 ± 1 V. The samples were exposed to the subsonic plasma jet for up to 1000 s. The sample surface temperature was 2000°C. The temperature was monitored with an optical pyrometer, which was focused onto the hot zone of the sample.  RESULTS AND DISCUSSION  Phase composition and morphology of materials forming  in  the Ta-C-Ir and Hf-C-Ir  systems.  According to X-ray diffraction data, the sample prepared by reacting a mixture of the tantalum carbide and iridium powders and silicon additions at 1600°C consisted of the TaC, TaIr3, and IrSi phases. Since the X-ray penetration depth in the system under consideration is approximately a few microns, this phase composition refers to the near-surface region of the sample. The intermetallic phase TaIr3 and iridium silicide were presumably formed by reactions (1) and (2), respectively:  TaC  +  3Ir  =  TaIr  +3  Ir  +  Si  =  IrSi.  C,  (1)  (2)  Figure 1 shows SEM images of the surface of the unoxidized sample. It is seen in Fig. 1a that the sample has an inhomogeneous surface covered with well-faceted crystals, which are connected by a continuous phase. The crystal size ranges up to 10 μm. Comparison of SEM/EDS and X-ray diffraction data indicates that the crystals consist of tantalum carbide, TaC. Note that the tantalum carbide powder used in our syntheses ranged in particle size from 200 to 300 nm. It is reasonable to assume that the preparation of the compact and subsequent heating to 1600°C followed by cooling lead to TaC recrystallization from a lowmelting-point eutectic of the Ir-Si system. Accumulations of TaC crystals have a golden color, typical of tantalum carbide. The other regions on the sample surface have a grey metallic luster. In the bulk of the compact, the intermetallic compound TaIr3 is a major phase. Distinctive morphological  features of  the material are the absence of open pores and strong cohesion of the grains (Fig. 1b).  In a similar way, we prepared samples from a mixture of HfC and iridium. The sample surface had a grey metallic luster. It also had regions differing in morphology (Fig. 2a). Large, well-faceted HfC single crystals ranging in size up to 10 μm were observed to form compact regions surrounded by a continuous light phase (Fig. 2b). SEM/EDS analysis of a crosssection of the sample in backscattered electron imaging mode showed that the sample contained phases differing in contrast, namely, the intermetallic phase HfIr3 and iridium silicide. It is reasonable to assume that reactions in the Hf-C-Ir system in the presence of a small silicon addition follow schemes similar to (1) and (2).  Thus, characteristic features of the microstructure formed as a result of the reaction of the metal carbide powders with iridium (in the presence of a small sili INORGANIC MATERIALS    Vol. 55    No. 3    2019  \\x0c', 'BEHAVIOR OF SOME REFRACTORY HAFNIUM  233  (a)  50 μm  (a)  25 μm  (b)  15 μm  Fig. 2. Electron micrographs of the material formed in the Hf-C-Ir system: (a) surface of the sample, (b) hafnium carbide crystals.   in a high-speed plasma f low showed that the two systems differed in behavior. The hafnium-containing sample tested at a temperature of 2000°C for 1000 s remained mechanically intact but became dull. The weight loss was ~0.3%. According to the X-ray diffraction data in Fig. 3a, the surface region of the oxidized sample consisted of monoclinic hafnium dioxide, hafnium silicate, and iridium. The formation of these phases can be represented by schemes (3)-(7):  HfIr  3  +  O  2  IrSi  2Ir  +  +  O  2  3O  2  HfO  +  2  SiO  +  2  3Ir,  Ir,  2IrO  3  ,  HfO  2  +  SiO  2  HfSiO ,  4  2HfC  +  3O  2  2HfO  +  2  2CO  .  (3)  (4)  (5)  (6)  (7)  (b)  100 μm  Fig. 1. Electron micrographs of the material formed in the Ta-C-Ir system: (a) survey micrograph of the sample surface, (b) cross-sectional morphology.   con addition) at a temperature of 1600°C are an essentially complete absence of pores throughout the sample, good cohesion of the grains of the different phases present, and the recrystallization of refractory tantalum or hafnium carbides, resulting in the formation of rather large single crystals. The forming microstructure is a consequence of the formation of an intermediate liquid phase, which adequately wets the surface of both the starting powders and reaction products and is favorable for sintering of all the components.  Phase composition and morphology of the materials after exposure to high-speed plasma f lows. Tests in air  INORGANIC MATERIALS    Vol. 55    No. 3    2019  → → → ↑ → → ↑ \\x0c', '234  y  t  i  s  n  e  t  n  I  Oxidized layer  BAKLANOVA et al.  (а)  HfSiO4 m-HfO2  (b)  δ-Ta2O5  (a)  25 μm  20  30  40  50 60 2θ, deg  70  80  90  Fig. 3. X-ray diffraction patterns of the samples oxidized in a plasma f low (surface temperature of 2000°C): (a) hafniumcontaining sample, (b) tantalum-containing sample.   More detailed information about the morphology and composition of  the oxidized  sample can be extracted  from cross-sectional SEM/EDS analysis data (Figs. 4, 5). It is seen from Fig. 4 that most of the sample was not oxidized and consisted of HfC, HfIr3, and iridium silicide grains very f irmly bonded to each other. No pores, through cracks across the sample, or mechanical damage was detected. The oxidized layer consisted of densely sintered, elongated hafnium dioxide and hafnium silicate crystals perpendicular to the surface of the oxidized sample. Note that the crystals sintered with each other formed layers parallel to each other and  the  sample  surface. SEM examination detected no glass-forming layer on the sample surface. The silicon dioxide forming by reaction (4) reacted with the hafnium dioxide to form crystalline hafnium silicate [reaction (6)], as evidenced by the X-ray diffraction data in Fig. 3a. On the whole, the laminated oxide layer readily peeled off from the surface of the unoxidized part, so it could not serve as a reliable barrier to oxygen diffusion. The workability of the hafnium-containing system under extreme conditions is due, f irst, to the pore-free microstructure of the material; second, to the very slow oxidation rate of iridium  (b)  100 μm  Fig. 4. Cross-sectional microstructure of the hafniumcontaining sample oxidized in a plasma f low: (a) oxidized layer, (b) cross-sectional SEM image of the sample in backscattered electron imaging mode.   even at 2000°C or higher temperatures; and, f inally, to the low vapor pressure of the forming solid hafnium dioxide.  Even though the tantalum system is similar in composition and microstructure to the hafnium system, its behavior during tests in a plasma f low differed from that of the latter system. When the sample surface temperature reached 2000°C, we observed the formation  INORGANIC MATERIALS    Vol. 55    No. 3    2019  \\x0c', 'BEHAVIOR OF SOME REFRACTORY HAFNIUM  235  6  7  5  4  3  2  1  100 μm  Spectrum  1  2  3  4  5  6  7  Hf  28.0  26.7  0  0  5.3  10.7  0  Atomic percent Ir  72.0  73.3  49.9  49.8  94.7  89.3  49.7  Si  0  0  50.1  50.2  0  0  50.3  Fig. 5. Micrograph of the hafnium-containing sample after exposure to a plasma f low and the elemental composition (EDS analysis data) of the sample in cross-sectional areas.   of a viscous liquid, which, driven by the plasma jet, shifted to the periphery of the sample. It is reasonable to expect that liquid formation was due to the melting of Ta2O5 (tm = 1877°C), a product of TaC and TaIr3 oxidation [13]. According to the X-ray diffraction data, after cooling the central part of the compact consisted of three phases: hexagonal Ta2O5, TaIr3, and TaC (Fig. 3b). According to the elemental analysis data, the upper layer consisted of not only tantalum and iridium but also silicon and oxygen. Therefore, it is reasonable to assume that the oxidized layer consisted of tantalum oxides and glassy silica. After cooling, the peripheral parts of the sample contained the triclinic and orthorhombic Ta2O5 phases, which was probably due to the difference in cooling rate between the hot, central part and the cooler, peripheral parts of the sample. The depth of the crater formed was about 0.4 mm and the weight loss after testing for 160 s was ~0.7%.  Like the hafnium-containing sample, the tantalum-containing sample remained intact, and no pores or cracks were detected (Fig. 6). Analysis of the crosssectional morphology of the sample after the test showed that no signif icant changes were produced, except for the formation of thin glassy surface layer. It is worth noting that at 2000°C Ta2O5 is in a liquid state, but its vapor pressure at this temperature is very low and its vaporization rate is as slow as 3.83 × 10-5 g/(cm2 s) [13]. Another inherent feature of the Ta2O5-SiO2 binary system resulting from oxidation is its immiscibility [14]. The presence of Ta2O5 in the glassy layer raises its liquidus temperature and viscosity, thereby impeding  INORGANIC MATERIALS    Vol. 55    No. 3    2019  oxygen diffusion across this layer. On the whole, the formation of liquid products in the Ta-C-Ir-Si system at the test temperature has a negative effect on the ablation resistance of the material, leads to the formation of a crater, and changes its surface topography.  CONCLUSIONS  Refractory  hafnium and  tantalum-containing materials have been prepared by reacting tantalum or hafnium carbide with iridium in the presence of a small amount of silicon. The synthesized materials consist of a mixture of phases: the intermetallic compound MIr3, recrystallized tantalum or hafnium carbide, and iridium silicide.  We have studied  the behavior of  the materials during a prolonged exposure to a high-speed plasma f low at a temperature of 2000°C. The results demonstrate that, owing to their special microstructure, the absence of pores, and the low oxidation rate of the iridium-containing components, both systems exhibit a good ablation resistance and that the hafnium system withstands a longer exposure.  ACKNOWLEDGMENTS  We are grateful to N.V. Bulina (Institute of Solid State Chemistry and Mechanochemistry, Siberian Branch, Russian Academy of Sciences) for collecting the X-ray diffraction patterns of the samples and to S.A. Terekhov (Central Research Institute for Special Machinery Joint Stock Company) for his assistance in    \\x0c', '236  BAKLANOVA et al.  performing the oxidation tests. This work was supported by  the Russian Science Foundation, grant no. 18-19-00075.  REFERENCES  (a)  1 mm  (b)  200 μm  (c)  10 μm  Fig. 6. SEM images of the tantalum-containing sample oxidized in a plasma f low: (a) general view of the sample surface, (b) oxidized layer, (c) cross-sectional SEM image of the sample in backscattered electron imaging mode.   2.  1. Ultra-High Temperature Ceramics: Materials for Extreme  Environment Applications, Fahrenholtz, W.G. et al., Eds., Hoboken: Wiley, 2014. Sevastyanov, V.G., Simonenko, E.P., Gordeev, A.N., Simonenko, N.P., Kolesnikov, A.F., Papynov, E.K., Shichalin, O.O., Avramenko, V.A., and Kuznetsov, N.T., HfB2-SiC (10-20 vol %) ceramic materials: manufacture and behavior under long-term exposure to dissociated air streams, Russ. J. Inorg. Chem., 2014, vol. 59, no. 12, pp. 1361-1382. 3. Parthasarathy, T.A., Rapp, R.A., Opeka, M., and Cinibulk, M.K., Modeling oxidation kinetics of SiC-containing refractory diborides, J. Am. Ceram. Soc., 2012, vol. 95, pp. 338-349. 4. Gild, J., Zhang, Y., Harrington, T., Jiang, S., Hu, T., Quinn, M.C., Mellor, W.M., Zhou, N., Vecchio, K., and Luo, J., High-entropy metal diborides: a new class of high-entropy materials and a new type of ultrahigh temperature ceramics, Sci. Rep., 2016, vol. 6, paper 37 946. 5. Wangping, W., Zhaofeng, C., Han, C., Liangbing, W., and Ying, Z., Tungsten and iridium multilayered structure by DGP as ablation-resistance coatings for graphite, Appl. Surf. Sci., 2011, vol. 257, no. 16, pp. 7295-7304.  6. Holleck, H., Binäre und Ternäre Karbidund Nitridsys teme der Übergangsmetalle, Berlin: Gebrüder Borntraeger, 1984.  9.  izdanie   Spravochnoe   7. Blagorodnye metally.   (Noble Metals: A Handbook), Savitskii, E.M., Ed., Moscow: Metallurgiya, 1984. 8. Yamabe-Mitarai, Y. and Murakami, H., Mechanical properties at 2223 K and oxidation behavior of Ir alloys, Intermetallics, 2014, vol. 48, pp. 86-92. Sha, J.B. and Yamabe-Mitarai, Y., Phase and microstructural evolution of Ir-Si binary alloys with fcc/silicide structure, Intermetallics, 2006, vol. 14, no. 6, pp. 672-684. 10. Lozanov, V.V., Baklanova, N.I., Bulina, N.V., and Titov, A.T., New ablation-resistant material candidate for hypersonic applications: synthesis, composition, and oxidation resistance of HfIr3-based solid solutions,  ACS Appl. Mater. Interfaces, 2018, vol. 10, no. 15,  pp. 13 062-13 072. 11. Criscione, J.M., Mercuri, R.A., Schram, E.P., Smith, A.W., and Volk, H.F., High temperature protective coatings  for graphite, part II, Technical Documentar y Report  ML-TDR-64-173, Air Force Materials Laboratory, 1964. 12. Strife, J.R., Smeggil, J.G., and Worrel, W.L., Reaction of iridium with metal carbides in the temperature range of 1923 to 2400 K, J. Am. Ceram. Soc., 1990, vol. 73, no. 4, pp. 838-845. 13. Kazenas, E. and Tsvetkov, Yu.,  (Vaporization of Oxides), Moscow: Nauka, 1997. 14. Reeve, D.A. and Bright, N.F.H., Phase relations in the system CaO-Ta2O5-SiO2, J. Am. Ceram. Soc., 1969, vol. 52, no. 8, pp. 405-409.  Isparenie oksidov  Translated by O. Tsarev  INORGANIC MATERIALS    Vol. 55    No. 3    2019  \\x0c']"
},{
  "_id": 21,
  "PDF": "Behavior of Ultra-High Temperature Ceramic Material HfB2–SiC–Y3Al5O12 under the Influence of Supersonic Dissociated Air Flow.pdf",
  "Text": "['ISSN 0036-0236, Russian Journal of Inorganic Chemistry, 2020, Vol. 65, No. 10, pp. 1596-1605. © Pleiades Publishing, Ltd., 2020. Russian Text © The Author(s), 2020, published in Zhurnal Neorganicheskoi Khimii, 2020, Vol. 65, No. 10, pp. 1397-1407.  INORGANIC MATERIALS AND NANOMATERIALS  Behavior of Ultra-High Temperature Ceramic Material HfB2-SiC-Y3Al5O12 under the Inf luence of Supersonic Dissociated Air Flow  E. P. Simonenkoa, *, N. P. Simonenkoa, A . N. Gordeevb, A . F. Kolesnikovb, A . S. Lysenkovc, I. A . Nagornova, d, V. N. Kurlove, A . E. Ershove, V. G. Sevast’yanova, and N. T. Kuznetsova  aKurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences, Moscow, 119991 Russia bIshlinsky Institute for Problems in Mechanics, Russian Academy of Sciences, Moscow, 119526 Russia cBaikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Moscow, 119334 Russia dMendeleev University of Chemical Technology of Russia, Moscow, 125047 Russia eInstitute of Solid State Physics, Russian Academy of Sciences, Chernogolovka, Moscow oblast, 142432 Russia *e-mail: ep_simonenko@mail.ru  Received April 6, 2020; revised May 29, 2020; accepted June 1, 2020  Abstract―Ultra-high-temperature ceramic (UHTC) materials with a density of 94.5 ± 1.3% have been manufactured by hot pressing a (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder, which has been prepared sol-gel method, at a moderate temperature of 1850°C (holding time, 30 min; pressure, 30 MPa). The oxidation resistance of this ceramic has been studied at elevated temperatures under the effect of a supersonic dissociated air f low (on a high-frequency induction plasmatron). The maximum surface temperature has been ~2550°C. An analysis of the kinetics of temperature changes depending on the heat load indicate that, when f ive percent of Y3Al5O12 by volume of HfB2-30 vol % SiC is introduced into the ceramics, the thermal conductivity of the material decreases. This is not critical from the point of view of the stability of the obtained sample to single temperature drops of ~700-1400°C in a few seconds with sharp heating and cooling of the sample. A decrease of ~35-40% in weight loss of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12) sample has been noted as the result of exposure as compared with unmodif ied analogs. Due to the presence of Y3Al5O12 in the initial ceramics, there is also a certain amount of the stabilized phase of cubic HfO2 in the oxidized layer in addition to the monoclinic.  Keywords: UHTC, HfB2, SiC, Y3Al5O12, thermochemical impact, high-enthalpy air f low, induction plasmatron  DOI: 10.1134/S0036023620100198  INTRODUCTION  In recent years, research has intensif ied in the f ield of creating ultra-high-temperature ceramic materials (UHTC) based on ZrB2/HfB2-SiC [1-10], which are supposed to be used at ultrahigh temperatures >2000- 2500°C; for example, to create the most thermally loaded parts of hypersonic aircraft, ceramic elements of promising propulsion systems [11, 12] or fuel cells for alternative energy [13, 14], in solar energy [15, 16], etc.  Special attention is currently being paid to f inding ways to improve the mechanical characteristics and the possibility of lowering the consolidation temperature of UHTC in order to avoid grain coarsening in this process. In particular, the introduction of various kinds of additives is discussed, which would make it possible to increase the strength and crack resistance of ceramics and their resistance to thermal shock without signif icant loss of oxidation resistance. The major ity of reports are devoted to the study of the effect of additives of carbon components of various nature [4, 9, 17-26]: f ibers, graphite or graphene plates, and nanotubes. A large number of studies are also associated with the introduction of refractory and superrefractory binary compounds into UHTC, such as carbides [27-32], silicides [33-36], and metal nitrides [36-40]. It is known that one of the most effective sintering additives that can signif icantly reduce the consolidation temperatures and, consequently, the risk of grain coarsening, primarily that of metal diborides, are metal oxides, the addition of which a priori should not increase the sensitivity of UHTC to oxidation. However, the number of works devoted to this topic is very small, probably due to concerns that the addition of a low-thermal-conductivity oxide component, which is distributed along the grain boundaries of HfB2 and SiC, causes a dramatic decrease in the thermal conductivity of the material as a whole.  1596  \\x0c', 'BEHAVIOR OF ULTRA-HIGH TEMPERATURE CERAMIC MATERIAL  1597  Nevertheless, there are a limited number of studies where  oxides  of  the Al2O3-Y2O3  system were attempted to be used as a sintering additive for ultrahigh-temperature ceramic materials based on zirconium diborides [41-45]. The composition of Y3Al5O12 is of particular interest; it possesses a high melting point, has no phase transformations in a wide temperature range [46, 47], and has a low capacity for oxygen diffusion [48, 49]. The insignif icant solubility of SiC in melts of the Al2O3-Y2O3 system [50] is used in the preparation of silicon carbide ceramics by liquidphase sintering [51, 52]. In all the available experimental studies on the manufacture of ceramic materials of the composition ZrB2-SiC-Al2O3-Y2O3, a positive role of the Y3Al5O12 additive on the compaction process was revealed, and in some works, an improvement in the resistance to oxidation in air at temperatures of 1600°C (the content of Y3Al5O12 was 10-40%) [41], 1700°C [44, 45], and short-term exposure to an oxygen-acetylene burner at temperatures of 2700- 2800°C [42, 44] was noted. Moreover, the possibility of stabilization of ZrO2 formed during the oxidation of ZrB2  in tetragonal YSZ was noted [44]. This can improve the adhesion of the oxidized UHTC area to the main body and reduce the chance of cracking and peeling. As far as we know, data on the preparation of ultra-high-temperature ceramic materials of the composition HfB2-SiC-Y3Al5O12 and their behavior in high-enthalpy air f lows have not reported. Additional arguments for initiating work on the introduction of Y3Al5O12 into the composition of the ZrB2/HfB2-SiC ceramics can be the results of a study of the vaporization of yttrium-aluminum garnet in the temperature range of 2200-2500°C [53]. The authors [53] found that at temperatures above 2200°C, more volatile aluminum oxide evaporates, while the yttrium oxide remaining in the system can stabilize the crystal lattice of resulting HfO2 in the tetragonal or cubic modif ication. At a lower temperature of the system, for example, in the volume of the oxidized part of the material under a layer of HfO2 with low-thermal conductivity, unevaporated Al2O3 can increase the viscosity of the protective glassy layer, reduce the rate of oxygen diffusion deep into the ceramic, and reduce the activity of silicon oxide during its evaporation. The aim of this work is to evaluate the behavior of the ultra-high-temperature ceramic material HfB2-30 vol % SiC modif ied with 5 vol % of highly dispersed Y3Al5O12 under the inf luence of a supersonic dissociated air f low.  EXPERIMENTAL  For the synthesis of aluminum and yttrium acetylacetonates required for the modif ication of HfB2- 30 vol % SiC, we used Y(NO3)3  · 6H2O  (99%, Khimmed), Al(NO3)3  · 9H2O  (99%, Khimmed),  acetylacetone C5H8O2 (>99.99%, EKOS-1), and 5% aqueous NH3 · H2O solution (specialty grade, Russian State Standard). n-Butyl alcohol C4H9OH (pure for analysis, Russian State Standard) was used as a solvent for the obtained chelate coordination compounds and a source of alkoxy groups in the subsequent synthesis of heteroligand complexes. Synthesis of composite powder (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 was carried out by the sol-gel technology [54] using metal alkoxoacetylacetonates as precursors (the methods for their preparation were described in detail [55-57]). The amount of the modifying additive Y3Al5O12 (5 vol %) was chosen as a compromise between the following factors: (i) the need to improve the ceramics compaction process through the use of the liquid-phase sintering method; (ii) the effect of increasing the resistance of the corresponding composite powder to oxidation in the temperature range of 20-1200°C when using the specif ied content of complex oxide, as shown [54]; (iii) the need to minimize the content of the UHTC oxide component, Y3Al5O12, which belongs to  low thermal conductivity substances (for polycrystalline Y3Al5O12, values from 3.2 (1273 K) to 8.7 W m-1 K-1 (296 K) [58] were obtained). Its addition can cause a signif icant decrease in the thermal conductivity of the ceramic material as a whole and  lead to disastrous destruction of the ceramic upon sharp heating or cooling. Production of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 ultra-high-temperature ceramic material by liquid-phase sintering was carried out using Thermal Technology Inc. hot-pressing (model HP20-3560-20) [29, 59-66]. Hot pressing was carried out at a temperature of 1850°C (holding time 30 min) with a heating rate of 10 deg/min; uniaxial pressure of 30 MPa was applied for the holding time after reaching the target temperature. The obtained sample was tested under rapid heating by a supersonic dissociated air f low on a VGU-4 100-kW high-frequency induction plasmatron. The sample f ixed in a water-cooled holder was introduced into the plasma jet at an anode power supply of the plasmatron (N) of 30 kW; N was increased to 60 kW with a step of 10 kW. The holding time at each stage at N = 30-60 kW was reduced in comparison with previous experiments [59, 60, 62-64] being equal to 2 min. Then the power was increased to 70 kW and the sample was held until the end of the experiment. The total exposure time was 2000 s (33 min 20 s). The study used a sonic nozzle with an outlet section diameter of 30 mm. The distance from the nozzle to the sample was 25 mm, the air f low rate was 3.6 g/s, and the pressure in the chamber was 13-14 hPa. To determine the average temperature of the sample surface, a Mikron M-770S pyrometer was used  in the spectral ratio pyrometer mode (temperature range 1000-3000°C, sighting area diameter ~5 mm); to register the temperature distribution over the sample surface, a Tandem VS-415U thermal  imager was used. Thermal  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020  \\x0c', '1598  m m  ,  e g a  k  n  i  r  h  S  SIMONENKO et al.  5  4  3  2  1  0  0  5  10 15 20 Exposure time, min  25  30  I  50  40  30  20  10  a  P  M  ,  e  r  u  s s  e  r  P  HfB2 SiC  10  20  30  40 2θ, deg  50  60  70  80  Fig. 2. X-ray powder diffraction pattern of the obtained HfB2-SiC-Y3Al5O12 sample.  Fig. 1. Shrinkage of the sample during hot pressing of the composite powder  (HfB2-30  vol % SiC)-5  vol % Y3Al5O12 at the stage of holding at 1850°C.  images were recorded at a wavelength of 0.9 μm at a specif ied value of the spectral emission coeff icient ε = 0.6. This value was selected on the basis of previous experiments [59, 60, 62-64, 67-71] according to the correspondence between  the  readings of  the pyrometer and thermal imager. Further, during the analysis of the thermal imager data, if necessary, the surface temperature values were corrected for real values of ε . The X-ray powder diffraction patterns of the surface of the samples before and after exposure to a supersonic dissociated air f low were recorded on a Bruker D8 Advance X-ray diffractometer (CuKα  radiation, 0.02° resolution with signal accumulation at a point for 0.3 s). X-ray powder diffraction was performed using the MATCH! software for phase identif ication from powder diffraction, version 3.8.0.137 (Crystal Impact, Germany), with the Crystallography Open Database (COD) integrated.  RESULTS AND DISCUSSION  Hot pressing of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder obtained by the method described [54] was carried out in graphite molds with an inner diameter of 15 mm using a small amount of boron nitride as a lubricant. Figure 1 shows the shrinkage at the f inal stage of holding the sample at a temperature of 1850°C. It is seen that an increase in pressure from 5 to 30 MPa, even without holding, leads to a sharp compaction of the material, probably due to the Al2O3-Y2O3-SiO2 liquid system existing on the grain surface [72-74]. The total holding time was 30 min; however, as can be seen in Fig. 1, no signif icant change in shrinkage occurs after the 20th minute. The material obtained by hot pressing has a density of 8.11 ± 0.11 g/cm3, which corresponds to 94.5 ±  1.3% of the theoretical value calculated by the additivity method (the density of HfB2, SiC, and Y3Al5O12 was taken equal to 11.2 g/cm3 [75], 3.2 g/cm3 [76], and 4.55 g/cm3 [77], respectively).  X-ray powder diffraction data (Fig. 2) indicate that no impurity crystalline phases (a by-product of carbothermal  synthesis HfC or a product of hafnium diboride HfO2 oxidation) were found in the sample. The X-ray diffraction patterns show extremely intense ref lections of the phase of hexagonal HfB2 [78] and cubic SiC [79]. The average crystallite size of silicon carbide  estimated  by  the Scherrer  formula was ~40 nm. Crystalline phases of  the Al2O3-Y2O3 or Al2O3-Y2O3-SiO2 systems were also not found in the composition; however, in the range 2 θ  ~ 8°-16° there is a diffuse halo, which may indicate the presence of a certain amount of an amorphous phase in the sample.  The most important characteristic of HfB2-SiC or ZrB2-SiC ultra-high-temperature ceramic materials is their good resistance to oxidation and the ability to withstand signif icant temperature drops (up to 1000- 1500°C in a few seconds) without disastrous destruction due to the high thermal conductivity of the materials (~55-80 W m-1 K-1 [80-85]), including elevated (>2000°C) temperatures. The introduction of even 5 vol % of Y3Al5O12 with low-thermal conductivity must lead not only to an increase in the resistance of the material to oxidation due to (i) more eff icient compaction during hot pressing, (ii) possible modif ication of the composition of silicate glass, and (iii) at least partial  stabilization of  the hafnium oxide phase formed during oxidation, but also to a decrease in the thermal conductivity of ceramics  in general. This explains the importance of studies of the behavior of the obtained ceramic material (HfB2-30 vol % SiC)- 5 vol % Y3Al5O12 under the conditions of relatively rapid heating by a supersonic dissociated air f low to temperatures >2000°C using an induction plasmatron that simulates aerodynamic heating as correctly as possible.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020      \\x0c', 'RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020  BEHAVIOR OF ULTRA-HIGH TEMPERATURE CERAMIC MATERIAL  1599  To make this true, the prepared cylindrical sample 15 mm in diameter and 4 mm thick was f ixed in a composite copper model using narrow strips of paper based on SiC whiskers in such a way as to prevent contact with the mandrel as much as possible. In order to minimize contact with the copper mandrel and reduce  heat transfer, the sample was mounted with a protrusion of 1 mm relative to the front surface of the model, as previously described in experiments with subsonic and supersonic high-enthalpy air f lows. Details of the process of installing a sample and assembling a watercooled model were reported [59, 60, 62, 63, 65].  The sample was introduced into a supersonic dissociated air f low at a steady-state power of the anode power supply of the plasmatron of 30 kW, which was then stepwise increased to 70 kW. The values of the exposure parameters and the average surface temperature of the samples determined using a spectral ratio pyrometer are given in Table 1 and Fig. 3.  At the f irst stage of exposure, there is the effect of a sharp heating of the surface due to the exothermic oxidation of hafnium diboride and silicon carbide followed by a slight (~30°C) temperature decrease (from 1490 to 1460°C) due to the formation of a protective glassy layer. With a subsequent increase in the anode power supply of the plasmatron and, accordingly, in the heat f lux, a stepwise increase in the surface temperature occurs. However, it should be noted that at N = 40-60 kW (q = 484-691 W cm-2), an increase in the average surface temperature is observed during holding at a constant power. In addition, the rate of temperature rise at each heating stage also tends to increase: it is 17 K/min at N = 40 kW, 21 K/min at N = 50 kW, and 34 K/min at N = 60 kW. At the beginning of the fourth stage (N = 60 kW), the temperature the critical values >1750-1850°C, after exceeded  which intensive evaporation from the surface of the protective glass layer usually occurs. In the middle of the eighth minute of exposure, bright points are visually observed on the surface overheated in comparison  Table 1. Change in the average surface temperature of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 sample in the vicinity of the critical point (spectral ratio pyrometer, T) depending on the exposure time and process parameters: the power of the anode supply (N) and pressure in the plasmatron chamber (P), as well as the corresponding values of the heat f lux1 (q)  1Heat f luxes to the water-cooled copper calorimeter were determined in separate experiments described [67].  Time, min  N, kW  P, hPa  q, W cm-2  Т, °C  0 → 2 → 4 → 6 →   2  30  13-14  363  1490 → 1590 → 1712 → 1846 →   1461   4  40  484   1623   6  50  598   1754   8  60  691   1914  8.5  70  779  ~2300  9  70  779  2520  10  70  779  2539  12  70  779  2540  16  70  779  2536  20  70  779  2539  25  70  779  2549  30  70  779  2553  33.33  70  779  2554  Fig. 3. Average temperature of the surface in the vicinity of the critical point of the obtained sample HfB2-SiC-Y3Al5O12 depending on the duration of the experiment and the parameters of action, namely the power of the anode supply N and the pressure in the plasma torch chamber P.  2400  T, °C 2600  2200  2000  1800  1600  1400  1200  1000  0  500  1000 Time, s  1500  2000  N, kW 80  70  60  50  40  30  P, hPa 30  25  20  15  10  T, °C  N, kW  P, hPa  \\x0c', '1600  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020  SIMONENKO et al.  with the average temperature. As a result of the transition to the maximum power of 70 kW, there is a very rapid (<1 min) increase in the surface temperature to ~2550°C, which does not change with small f luctuations until the end of the experiment.  Analysis of the thermal images of the sample surface (Fig. 4) also indicates that as the power of the anode supply increased, the temperature of the sample surface increased. Moreover, until about the middle of the eighth minute of the test, the temperature distribution over the sample surface was relatively uniform, with the exception of a thin ring along the edge of the sample ~0.5-1 mm thick. However, after overheated microregions appeared, a  temperature gradient  is formed along the radius: in the central region, the temperature becomes higher (by ~70-80°C) than at the periphery.  The temperature difference on the sample surface increases from 150 to 520°C in a fairly narrow time interval (530-539 s of the experiment). The number of  overheated surface areas grows rapidly during this time; they begin to join and by 555 s of exposure they f ill almost the entire area. In this case, the average surface temperature is ~2420-2500°C. The kinetics of surface temperature changes at specif ic points is illustrated in Fig. 5. In ~10 min after the beginning of the exposure, the surface temperature stabilizes at 2540- 2550°C. Nevertheless,  throughout  the experiment, there are small elevation differences on the surface, which are probably associated with relief inhomogeneities. The  latter are traces of bubbles that burst during evaporation from the surface of silicate glass, which during their formation displaced refractory particles of hafnium oxide, as well as traces of erosion of the oxidized layer. The equalization of temperatures at the periphery of the sample and in the center (Fig. 5) after the 25th minute may indicate the removal of convex formations appeared during the sharp evaporation of the vitreous layer, which occurred most intensively in the central region.  Fig. 4. Thermal images illustrating temperature distribution over the surface of the HfB2-SiC-Y3Al5O12 sample at various stages of exposure to a supersonic f low of dissociated air, as well as along the sample diameter indicated by a straight line.  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  621 s  2600 °С 2200  1800  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  727 s  2600 °С 2200  1800  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  1124 s  2600 °С 2200  1800  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  1685 s  2600 °С 2200  1800  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  2000 s  2600 °С 2200  1800  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  555 s  2500 °С 2100  1600  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  544 s  2500 °С 2100  1600  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  539 s  2400 °С 2000  1600  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  536 s  2100 °С  1700  1300  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  530 s  1900 °С  1500  1100  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  70 s  1500 °С 1300  1100  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  150 s  1500 °С 1300  1100  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  250 s  1600 °С  1300  1100  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  350 s  1700 °С  1400  1100  2500  2000  1500  1000 -20  0 R, mm  20  T  ,  °  C  475 s  1800 °С  1500  1200                                \\x0c', 'BEHAVIOR OF ULTRA-HIGH TEMPERATURE CERAMIC MATERIAL  1601  2500  445 s  T3  T1  1800 °С 1500  2000  T2  С  °  T1  1200  r  u  ,  e  t  a  r  e  p  m  e  T  1500  1000  0  T2  T3  T3  2600 °С 2000  T2  T1  2000 s  2600 °С 2100  1700  T3  T2  T1  570 s  T3  T2  T1  545 s  2600 °С 2100  1700  530 s  T3  T2  T1  2000 °С 1700  1400  500  1000 Time, s  1500  2000  Fig. 5. Changes in the temperature of individual sections of the surface of the HfB2-SiC-Y3Al5O12 sample in the course of exposure to a supersonic f low of dissociated air.  Thus, it can be stated that, despite the presence of 5 vol % Y3Al5O12 with a low thermal conductivity, the obtained sample withstood the effect of a supersonic dissociated air f low like unmodif ied samples of the composition HfB2-30 vol % SiC [62, 63]: The destruction of the sample, cracking or delamination of the oxidized part of the material was observed neither during heating nor during sharp cooling. Higher (by ~80-150°C) surface temperatures of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 sample at the  initial stages of exposure  in comparison with unmodif ied materials [62, 63] can be explained by its lower thermal conductivity, a decrease in which, obviously, is not critical for resistance to a single thermal shock. A sharp rise in temperature (the so-called temperature jump [86]) began for the obtained material at the same heat f lux (q = 779 W cm-2) as for materials HfB2-30 vol % SiC; however, the initial temperature of such a jump was higher: ~1900°C compared to ~1800°C for the latter [62, 63].  The weight loss of the sample after exposure is 1.7%, which  is  signif icantly  less  than  the values obtained earlier for materials HfB2-30 vol % SiC (2.6-3.0%) [62, 63] taking into account the fact that the surface temperature in all cases was about 2550- 2600°C for almost the same time (24-25 min). This can be explained by the inf luence on the properties of the protective vitreous layer of such introduced glassforming  refractory components as aluminum and yttrium oxides, which are often purposefully added to the composition of glasses to increase their refractoriness.  X-ray powder diffraction analysis of the oxidized surface (Fig. 6) indicates that the main crystalline phase among the oxidation products is monoclinic HfO2 [87]. This is typical for samples exposed to a suff iciently long exposure to temperatures of ~2500- 2600°C. No impurity of hafnium silicate HfSiO4 was found. However, in the region 2 θ  = 30°-31° there is a broadened ref lection of relatively low intensity, which can be attributed to both the f luorite (cF12) [88] and orthorhombic (oP16) [89] phases of stabilized HfO2. The content of this more ordered hafnium oxide phase is ~3-6%, which corresponds to a small amount of the introduced sintering additive Y3Al5O12.  CONCLUSIONS  Ultra-high-temperature ceramic materials with a density of 94.5 ± 1.3% were prepared by hot pressing of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder obtained by the sol-gel method at a moderate temperature of 1850°C (holding time, 30 min; applied pressure, 30 MPa). X-ray powder diffraction showed that only HfB2 and SiC crystalline phases were present in the composite ; the oxide component was in an amorphous state, probably within the Al2O3-Y2O3-SiO2 system.  To study the resistance to oxidation at elevated temperatures under the action of atomic oxygen, the sample was introduced into a supersonic jet of dissociated air (a high-frequency  induction plasmatron). Heat f luxes directed to the water-cooled copper calorimeter in the experiment varied stepwise from 363 to  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020    \\x0c', '1602  SIMONENKO et al.  I  10  20  30  experiment  2θ, deg 40  50  I  60  70  80  Experiment HfO2, mP12 HfO2, cF12 HfO2, oP16  28  30 2θ, deg  32  HfO2, mP12  HfO2, cF12  HfO2, oP16  10  20  30  40 2θ, deg  50  60  70  80  Fig. 6. X-ray powder diffraction pattern of the surface of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 sample oxidized by a supersonic f low of dissociated air (black) and reference X-ray diffraction patterns: monoclinic hafnium oxide (red) [87], cubic (cF12, blue) [88], and orthorhombic phases (oP16, green) [89] of HfO2.  779 W cm-2 as the anode power supply of the plasmatron increased (the pressure in the chamber was 13- 14 hPa).  It was found that, starting from the second stage of heating (q = 484 W cm-2), when the heat f lux is f ixed, there is a tendency to a gradual increase in the surface temperature, and as the power of the anode power supply of the plasmatron increases, the rate of temperature change also increases. At a surface temperature of ~1900°C, at which local overheated regions were appeared, when passing to the maximum power N (q = 779 W cm-2), a sharp (<1 min) increase in the surface temperature occurs up to ~2400-2500°C. In this case, there is an increase in the number and area of high-temperature microregions on the surface until  the temperature stabilizes at 2540-2550°C 10 minutes after the start of heating.  Analysis of the kinetics of temperature changes depending on the heat load may indicate that the thermal conductivity of the material in general decreases as a result of the introduction of 5 vol % Y3Al5O12 with low-thermal conductivity into the HfB2-30 vol % SiC ultra-high-temperature ceramics. However, this was found to be not critical from the point of view of the stability of the obtained sample to single temperature drops of ~700-1400°C in a few seconds: neither heating when injecting dissociated air into a high-enthalpy jet, nor cooling when heating was turned off led to the destruction of the sample.  It was shown that the weight loss after holding the obtained (HfB2-30 vol % SiC)-5 vol % Y3Al5O12  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020  \\x0c', 'BEHAVIOR OF ULTRA-HIGH TEMPERATURE CERAMIC MATERIAL  1603  sample at a surface temperature of ~2550°C for 25 min (with a total test time of >33 min) was found to be ~35-40% less than it was determined [62, 63] for materials HfB2-30 vol % SiC under a similar effect being 1.7%. This may be the result of a modif ication of the composition and, therefore, the viscosity of the protective vitreous layer due to the introduction of Y3Al5O12 into the composite. X-ray powder diffraction conf irmed the assumption of a a possible partial stabilization of the HfO2 formed during the oxidation of hafnium diboride: in the composition of the surface oxidized layer, in addition to monoclinic hafnium oxide, HfO2 stabilized in the orthorhombic or, more likely, cubic modif ication was present. The content of this phase was ~3-6%. For a correct assessment of the effect on adhesion of its presence in the oxidized near-surface layer, additional experiments are required. Thus, the introduction of 5 vol % Y3Al5O12 in the form of a thin f ilm on the surface of HfB2-SiC powder particles made it possible to improve the hot pressing of an ultra-high-temperature ceramic material, as well as to increase its resistance to the action of a supersonic f low of dissociated air at a surface temperature of ~2550°C.  FUNDING  The study was supported by the Russian Science Foundation (project no. 17-73-20181).  CONFLICT OF INTEREST  The authors declare that they have no conf lict of interest.  REFERENCES  1. E. P. Simonenko, D. V. Sevast’yanov, N. P. Simonenko, et al., Russ. J. Inorg. Chem. 58, 1669 (2013).  https://doi.org/10.1134/S0036023613140039  2. C. M. Carney, in Comprehensive Composite Materials II  (Elsevier, 2018), p. 281.  https://doi.org/10.1016/B978-0-12-803581-8.09996-3 3. R. Savino, L. Criscuolo, G. D. 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Chae, S.-M. Lee, et al., J. Korean Ceram. Soc. 50, 410 (2013).  https://doi.org/10.4191/kcers.2013.50.6.410 84. S. Kim, J.-M. Chae, S.-M. Lee, et al., Ceram. Int. 40, 3477 (2014).  https://doi.org/10.1016/j.ceramint.2013.09.082 85. D. Debnath, S. Chakraborty, A. R. Mallick, et al., Adv. Appl. Ceram. 114, 45 (2015).  https://doi.org/10.1179/1743676114Y.0000000189 J. Marschall, D. Pejakovic, W. G. Fahrenholtz, et al., J. Thermophys. Heat Transfer 26, 559 (2012).  https://doi.org/10.2514/1.T3798 87. R. P. Haggerty, P. Sarin, Z. D. Apostolov, et al., J. Am. Ceram. Soc. 97, 2213 (2014).  https://doi.org/10.1111/jace.12975 88. D.-J. Kim, S.-H. Hyun, S.-G. Kim, and M. Yashima, J. Am. Ceram. Soc. 77, 597 (1994).  https://doi.org/10.1111/j.1151-2916.1994.tb07035.x J. Tang, M. Kai, Y. Kobayashi, et al., Geophysical  Monograph, Vol. 101: Properties of Earth and Planetary Materials at High Pressure and Temperature (Wiley,  1998), p. 401.  https://doi.org/10.1029/GM101p0 401  86.  89.  Translated by V. Avdeeva  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 10    2020  \\x0c']"
},{
  "_id": 22,
  "PDF": "Behavior of Ultrahigh-Temperature ZrB2-Based Ceramics in Oxidation.pdf",
  "Text": "['DOI 10.1007/s11106-018-9930-z  Powder Metallurgy and Metal Ceramics, Vol. 56, Nos. 9-10, January, 2018 (Russian Original Vol. 56, Nos. 9-10, Sept.-Oct., 2017)   BEHAVIOR OF ULTRAHIGH-TEMPERATURE   ZrB2-BASED CERAMICS IN OXIDATION   O. N. Grigoriev,1 I. P. Neshpor,1 T. V. Mosina,1,5 V. B. Vinokurov,1  A. V. Koroteev,1 O. V. Buriachek,2 D. V. Vedel,3  A. N. Stepanchuk,3 and L. Silvestroni4   UDC 663.3.7:620.193   Ultrahigh-temperature ZrB2-based ceramics with different sintering additions was developed for   extreme conditions. Its strength characteristics, phase composition, and structure were examined.   The ceramics was oxidized in air at 1250 and 1550°C. In addition, the most stable composites were   subjected to temperature cycling in a flow of aviation fuel combustion products in a temperature   range of 1400-1500°C. All materials show high oxidation resistance. The method used to produce   samples influences their oxidation behavior:  materials produced by vacuum hot pressing show  higher oxidation resistance than those produced by hot pressing in a CO-CO2 atmosphere, probably  because of their higher final density. The best results were obtained when ZrB2 sintering was  combined with introduction of MoSi2 and CrB2.   Keywords: oxidation, ultrahigh-temperature ceramics, zirconium diboride, temperature cycling test.   INTRODUCTION   Ultrahigh-temperature ceramics (UHTC) is used for thermal protection of hypersonic aerospace vehicles or   reusable space vehicles and for the manufacture of special components for gas turbine engines, furnaces, refractory   crucibles, etc. Sharp front edges and nose spinners of hypersonic space vehicles operate at high temperatures and in   thermal cycling conditions in an oxidizing atmosphere. The materials must possess high oxidation resistance,   adequate thermal shock properties, and low creep. Modern thermal protection materials, which are limited to SiC  and Si3N4 ceramics, oxide ceramics, and C/C composites, have high oxidation resistance (up to 1600°C), but low  long-term resistance. Ultrahigh-temperature ZrB2 materials are most promising for extreme applications as they   combine high melting temperatures with thermal conductivity and with high erosion and corrosion resistance [1-9].  To examine the effect of various sintering additives on the corrosion resistance of ZrB2 UHTC produced by hot   pressing and vacuum hot pressing (HP and VHP), the materials were oxidized in air at 1550°C.    Our objective is to examine the corrosion behavior of zirconium diboride ceramics of various compositions   at different temperatures, analyze how the production technique and phase composition influence resistance in an   1Frantsevich Institute for Problems of Materials Science, National Academy of Sciences of Ukraine, Kyiv,   Ukraine.   2Ivchenko-Progress Machine-Building Design Office, Zaporizhzhia, Ukraine.   3National Technical   University “Igor Sikorsky Kyiv Polytechnic Institute,” Kyiv, Ukraine. 4Institute of Science and Technology for   Ceramics, Faenza, Italy.   5To whom correspondence should be addressed; e-mail: mosinatv@ukr.net.   Translated from Poroshkovaya Metallurgiya, Vol. 56, Nos. 9-10 (517), pp. 110-119, 2017. Original article   submitted December 7, 2016.       1068-1302/18/0910-0573 \\uf0d32018 Springer Science+Business Media, LLC                            573\\xa0  \\x0c', '574   oxidizing atmosphere, and test the materials in thermal cycling conditions to establish their residual mechanical   properties in extreme applications.   EXPERIMENTAL PROCEDURE   The ZrB2-based composites were produced by HP without a protective atmosphere and by VHP. The   characteristics of starting powders are summarized in Table 1.   Oxidation tests were performed in air in nonisothermal conditions at a heating rage of 3-4 °C/min at 1250   and 1550°C for 30 min with slow furnace cooling. The samples were oxidized in a VMK 1600 furnace (Linn High   Term).   The density of   the   samples was measured with   the Archimedes methods   and   evaluated   from   microphotographs (scanning electron microscopy (SEM)) processed with the Image Pro software.   Before the tests, all samples were cleaned in solvents, weighed, and measured (length, thickness, and   width). The weight increment after oxidation was estimated by calculating the ratio between the weight increment,   resulting from the formation of oxides, and the specific surface area before the experiment.  In addition, the behavior of ZrB2-MoSi2 ceramics in thermal cycling in a flow of flamethrower combustion   products in the 1400-1500°C range was examined using a test bench at the Progress Design Office.    The morphology of the powders, microstructure of the sintered ceramics, and microstructural changes   induced by oxidation were evaluated using X-ray diffraction, optical microscopy, and   scanning electron   microscopy.   TABLE 1. Characteristics of Starting Powders for UHTC   Powder   Supplier   Structure   Powder size, µm   Purity, % Admixtures, wt.%  ZrB2   Beijing Institute of Aeronautic  Materials, China  Same   Hexagonal   1.0   99.0   0.8 O, 0.3 C   SiC  MoSi2  ZrB2   Hexagonal   0.5   99.9   0.1 O   Same   Tetragonal   1.0-3.0   99.9   0.1 O   Frantsevich Institute for Problems of  Materials Science, Ukraine    Hexagonal   10   98   1.5 O, 0.4 C   MoSi2  SiC  CrB2   Same   Tetragonal   5   98.5   1.2 O, 0.3 Fe   Zaporizhzhia Abrasive Plant, Ukraine  Donetsk Chemical Reagent Plant,  Ukraine    Hexagonal  Hexagonal   5-7  7.9   99  99.1   0.9 O, 0.2 Fe  0.3 O, 0.09 C, 0.2  Fe   TABLE 2. Composition, Sintering Parameters, and Density of ZrB2 Ceramics and Phases Formed   after Oxidation of VHP Ceramics    Ceramics, vol.%   Sintering parameters*   Density, g/cm3   XRD   T, °C   \\uf074, min   measured   theoretical   relative, %   ZrB2 -15MoSi2  ZrB2 -3SiC-5MoSi2  ZrB2 -15SiC-5MoSi2  ZrB2 -15SiC  ZrB2 -15SiC-5CrB2  ZrB2 -15MoSi2-5CrB2   2010     9   5.38   6.19   86.8   ZrB2,** SiC  ZrB2, SiC  ZrB2,** SiC  ZrB2, SiC  ZrB2,** SiC  ZrB2,** MoSi2, MoB   1870   15   5.64   6.09   92.7   1870   15   5.31   5.73   92.6   2075   15   5.55   5.66   99.9   1870   20   5.50   5.68   97.1   1740   20   5.79   6.14   93.9   *Sintering pressure 48 MPa. **Solid solution.       \\x0c', 'TABLE 3. Properties of Ceramic Composites at Room Temperature    Ceramics, vol.%   Vickers hardness   (1 kg), GPa   Fracture toughness  (10 kg), MPa \\uf0d7 m0.5   Bending strength, MPa    ZrB2 -15MoSi2  ZrB2 -3SiC-5MoSi2  ZrB2 -15SiC-5MoSi2  ZrB2 -15SiC-5CrB2  ZrB2 -15MoSi2-5CrB2  ZrB2 -15MoSi2   21.3 ± 0.5   22.4 ± 0.6   17.1 ± 0.4   19.6 ± 0.5   22.4 ± 0.6   19.1 ± 0.4   3.58 ± 0.11   -   5.32 ± 0.3   2.00 ± 0.20   5.79 ± 0.24   4.09 ± 0.14   490 ± 50   420 ± 65   500 ± 35   495 ± 40   370 ± 30   365 ± 55   The composition, sintering parameters, and density of the ceramics and phases that formed after oxidation  are shown in Table 2. The mechanical properties of the ceramics produced by HP in a CO/CO2 atmosphere are   provided in Table 3.   a   c   b   d   e   Fig. 1. Surface microstructure of ceramics produced by HP in a CO-CO2 atmosphere after oxidation  at 1250°C for 30 min, vol.%: a) ZrB2-15 MoSi2; b) ZrB2-3 SiC-5 MoSi2; c) ZrB2-15 SiC-15 MoSi2;   d) ZrB2-15 SiC-5 CrB2; e) ZrB2-15 MoSi2-5 CrB2   575                   \\x0c', 'OXIDATION BEHAVIOR   Oxidation of ZrB2 Ceramics. Hot-pressed ceramic samples (Table 2) were oxidized in air in nonisothermal   conditions at 1250 and 1550°C for 30 min with slow cooling in the furnace.   Visual examination of the oxidized samples allows us to obtain information on the behavior of ceramics in   oxidation. Externally, color of the ceramics oxidized at 1250°C for 30 min varies from light gray to dark gray.   Oxidation at 1550°C leads to the formation of a whitish-yellowish film on the samples.   The samples oxidized at 1250°C are shown in Figs. 1 and 2 and those oxidized at 1550°C in Figs. 3 and 4.   The weight increment and phase composition of oxidation products of the ceramics produced by hot pressing and   vacuum hot pressing are presented in Table 4. Considering the specific weight increment, note that all composites  have high oxidation resistance and introduction of 3-15 vol.% SiC or CrB2 further improves the corrosion resistance of    a   c   b   d   Fig. 2. Surface microstructure of ceramics produced by VHP after oxidation at 1250°C for 30 min, vol.%:  a) ZrB2-15 MoSi2; b) ZrB2-15 SiC-15 MoSi2; c) ZrB2-15 SiC-5 CrB2; d) ZrB2-15 MoSi2-5 CrB2   TABLE 4. Specific Weight Increment and Phase Composition of Ceramic Oxidation Products    Production  method   Material, vol.%   Specific weight  increment, mg/cm2  XRD   1250\\uf0b0C   1550\\uf0b0C   1250\\uf0b0C   1550\\uf0b0C   HP   VHP   ZrB2-15MoSi2  ZrB2-3SiC-5MoSi2  ZrB2-15SiC-5MoSi2  ZrB2-15SiC-5CrB2  ZrB2-15MoSi2-5CrB2  ZrB2 -15MoSi2  ZrB2 -15SiC-5MoSi2  ZrB2 -15SiC-5CrB2   ZrB2 -15MoSi2-5CrB2   1.4   5.7   1.0   0. 2   0.8   -1.3   0. 6   -1.5   0.6   1.6   2.5   3.9   1.7   1.3   -1.8   1.8   -1.8   3.5   ZrO2, SiO2  ZrO2  ZrO2, SiO2  SiC, ZrSiO4, ZrO2, ZrB2  ZrSiO4, SiO2, ZrO2  ZrO2, SiO2, ZrB2  ZrB2, ZrO2, SiC   ZrB2, ZrO2, SiO2, ZrSiO4, SiC  SiO2, ZrSiO4, ZrO2  ZrO2, ZrSiO4  ZrSiO4, SiO2, ZrO2, SiC  ZrSiO4, ZrO2, SiO2  ZrSiO4, ZrO2, SiO2, CrMoO4  ZrB2, ZrO2, SiO2, ZrSiO4  ZrO2, ZrSiO4, SiO2, ZrB2, SiC ZrO2, SiO2, ZrSiO4   ZrB2, SiO2, ZrO2   ZrB2, ZrO2, SiO2, ZrSiO4   576                 \\x0c', 'a   c   b   d   Fig. 3. Surface microstructure of ceramics produced by HP in a CO-CO2 atmosphere after oxidation  at 1550°C for 30 min, vol.%: a) ZrB2-15 MoSi2; b) ZrB2-3 SiC-5 MoSi2; c) ZrB2-15 SiC-15 MoSi2;   d) ZrB2-15 SiC-5 CrB2; e) ZrB2-15 MoSi2-5 CrB2   e   a   c   b   d   Fig. 4. Surface microstructure of ceramics produced by VHP after oxidation at 1550°C for 30 min, vol.%:  a) ZrB2-15 MoSi2; b) ZrB2-15 SiC-15 MoSi2; c) ZrB2-15 SiC-5 CrB2; d) ZrB2-15 MoSi2-5 CrB2   577                         \\x0c', 'Fig. 5. XRD of ZrB2-15% MoSi2-5% CrB2 ceramics oxidized at 1500°C   a   b   c   Fig. 6. SEM analysis of cross-sections of oxidized samples, vol.%: ZrB2-15 MoSi2 (a), ZrB2-15  MoSi2-5% CrB2 (b), and ZrB2-15 SiC-5 CrB2 (c)   578             \\x0c', 'Fig. 7. Thermal cycling test in flamethrower blast   materials with the formation of a protective film consisting of high-temperature oxides and zirconium. The weight   decrement can be explained by evaporation of boron oxide or loss of oxide film debris.   X-ray diffraction (XRD) (Fig. 5) revealed a halo on the ceramics, which is associated with the formation of  a SiO2-B2O3-based amorphous phase. A dense glassy film forms on all ceramics oxidized at 1250°C. At 1550°C,   the surface morphology varies considerably and is generally rough and uneven with yellowish-white inclusions of  oxidation products, except for the samples containing an addition of CrB2. The oxidation resistance of the ceramics  is relatively the same at low temperatures, while addition of CrB2 significantly improves its high-temperature   corrosion resistance at 1500°C.  X-ray diffraction shows that the ZrB2 phase is retained and revealed in all samples, even those oxidized at   1550°C. The difference in the behavior of samples produced by HP and VHP is associated with high density of the  latter. In most cases, zircon (ZrSiO4) is formed besides ZrO2 and SiO2, especially at 1550°C. Zircon is a hightemperature phase  that ensures corrosion resistance of  the composite  to 1550-1600°C [10]. When CrB2  is  introduced into two-component composites, ZrSiO4 shows up already at 1250°C.   The cross-section of the oxidation zone was examined with SEM. Microphotographs of the oxidized  samples are shown in Fig. 6. The oxide layer of materials has a thickness of 50-100 \\uf06dm and is quite dense. It is a   multilayer structure: the outer scale layer consists of silicon oxide, a layer that follows contains silicon and  zirconium oxides and zircon, and the next layer is represented by the starting ZrB2, MoSi2, and CrB2 phases and  Zr(Mo)B2 solid solution.   Behavior of ZrB2-MoSi2 Ceramics in Thermal Cycling. The ZrB2-15% MoSi2 ceramic materials were   tested by thermal cycling in a flow of aviation fuel combustion products, in a flamethrower blast, at 1400-1500°C.   The tests were performed using the test bench at the Progress Design Office (Fig. 7).   The samples after tests with a different number of cycles are shown in Fig. 8. The surface defects that   formed during abrasive processing are healed when the ceramics oxidize over 1000-3000 cycles. This is confirmed   by increase in bending strength of the ceramics from 350 to 450-500 MPa.   a   b   Fig. 8. Samples after tests with 1850 (a) and 3975 (b) cycles: \\uf044T = 1400°C   Fig. 9. Residual strength of ceramics after thermal cycling: \\uf044T = 1400°C, N is the number of cycles   579               \\x0c', 'The ceramics retain high strength in further tests up to 6000 cycles (Fig. 9). This means that strength does   not deteriorate in thermal cycling tests with 6000 cycles. Hence, there are prospects for application of the studied   ceramics in high-temperature operating conditions.   The ceramic samples were tested for bending at room temperature and 1400°C. It turned out that the  samples had greater bending strength at higher temperatures\\uf0be500-600 MPa.   CONCLUSIONS   The high-temperature behavior and oxidation resistance of ZrB2 ceramic composites have been studied. All   materials possess high oxidation resistance. The production method   influences   the behavior of samples:   the   materials produced by hot pressing have higher resistance than the materials produced by vacuum hot pressing in a  CO-CO2 atmosphere.   The surface defects on the ceramics are healed in the oxidation process during thermal cycling tests with  \\uf044T =1400°C over 1000-3000 cycles. The strength of the ceramics increases from 450 MPa to 500-550 MPa. High   strength is retained in subsequent thermal cycling tests with up to 6000 cycles.    The tests performed in 6000 cycles indicate that the strength does not degrade; this is a favorable factor for   high-temperature applications of this ceramics.   It is established that the bending strength at 1400°C is 500-600 MPa, which is higher than at room   temperature.   1.   2.   3.   4.   5.   6.   7.   8.   9.   REFERENCES   Shu-Qi Guo, “Densification of ZrB2-based composites and their mechanical and physical properties: A   review,” J. Eur. Ceram. Soc., 29, 995-1011 (2009).  F. Monteverde and A. Bellosi, “Oxidation of ZrB2-based ceramics in dry air,” J. Electrochem. Soc., 150,   552-559 (2003).   Hu Ping, Zhang Xing-Hong, and Han Jie-Cai, “Effect of various additives on the oxidation behavior of  ZrB2-based ultrahigh-temperature ceramics at 1800°C,” J. Am. Ceram. Soc., 93, Issue 2, 345-349 (2010).  O. N. Grigoriev, B. A. Galanov, V. A. Lavrenko, et al., “Oxidation of ZrB2-SiC-ZrSi2 ceramics   in   oxygen,” J. Eur. Ceram. Soc., 30, 2397-2405 (2010).   J. F. Justin and A. Jankowiak, “Ultrahigh-temperature ceramics: densification, properties and   thermal   stability,” J. Aerosp. Lab., 3, 1-10 (2011).  D. Sciti, R. Savino, and L. Silvestroni, “Aerothermal behavior of a SiC fiber-reinforced ZrB2 sharp   component in supersonic regime,” J. Eur. Ceram. Soc., 32, 1837-1845 (2012).   V. O. Lavrenko, A. D. Panasyuk, O. M. Grigoriev, et al., “High-temperature (to 1600°C) oxidation of  ZrB2-MoSi2 ceramics in air,” Powder Metall. Met. Ceram., 51, No. 1-2, 102-107 (2012).   O. N. Grigoriev, G. A. Frolov, U. I. Evdokimenko, et al., “Ultrahigh temperature ceramics behavior under   the impact of concentrated solar radiation, oxidation and erosion in gas flows,” in: Space Investigations in   Ukraine 2014, Akademperiodyka, Kyiv (2014), pp. 126-132.   T. V. Mosina,   I. P. Neshpor, O. M. Grigoriev, et al., “Corrosion   resistance of ultrahigh-temperature   zirconium boride ceramics under concentrated solar radiation,” Powder Metall. Met. Ceram., 54, No. 3-4,   10.   189-193 (2015).  A. Kaiser, M. Lobert, and R. Telle, “Thermal stability of zircon (ZrSiO4),” J. Eur. Ceram. Soc., 28, 2199-  2211 (2008).   580       \\x0c']"
},{
  "_id": 23,
  "PDF": "Beneficial effects of an ultra-fine α-SiC incorporation on the sinterability and mechanical properties of ZrB2.pdf",
  "Text": "['Appl. Phys. A 82, 329 -337 (2006)  DOI: 10.1007/s00339-005-3327-9  Applied Physics A  Materials Science & Processing  f. monteverde  Beneficial effects of an ultra-fine α-SiC incorporation on the sinterability and mechanical properties of ZrB2  ISTEC-CNR, Via Granarolo 64, 48018 Faenza, Italy  Received: 9 June 2005/Accepted: 22 June 2005 Published online: 25 August 2005 • © Springer-Verlag 2005     √  ABSTRACT A fully dense ZrB2 ceramic containing 10 vol. % ultrafine α-SiC particulate was successfully hot pressed at 1900 ZrB2 grains (average size ≈ 3 µm) and SiC particles dispersed C for 20 min and 40 -50 MPa of applied pressure. Faceted regularly characterized the base material. No extra secondary phases were found. The introduction of the ultrafine α-SiC particulate was recognized as the key factor that enabled both the control of the diboride grain growth and the achievement of ing combination of data: 4.8 ± 0.2 MPa full density. The mechanical properties offered an interestm fracture toughness, 507 ± 4 GPa Young’s modulus, 0 .12 Poisson’s ratio, and 835 ± 35 MPa flexural strength at C (in air) provided values of 300 ± room temperature. The flexural strength measured at 1500 35 MPa. The incorporated ultrafine α-SiC particulate was fundamental, sinterability apart, to enhancing the strength and oxidation resistance of ZrB2 . The latter property was tested at 1450 C for 20 h in flowing dry air. In such oxidizing conditions, the formation of a thin external borosilicate glassy coating supplied partial protection for the faces of the material exposed to the hot environment. The oxidation attack penetrated into the material’s bulk and created a 200µm-thick zirconia scale. The SiC particulate included in the oxide scale, lost by active oxidation, left carbon-based inclusions in the formerly occupied sites.        PACS 81.05.Je; 81.20.Ev; 81.70.Bt  1  Introduction     IV-V-group transition-metal diborides and carbides possess melting temperatures among the highest known, i.e. above 3000 C. This peculiarity makes them potential candidates for use at very high temperatures. Zirconium diboride (ZrB2 ), for instance, has recently been considered for use in the manufacture of electrodes or crucibles for moltenmetal contact, to name but a few examples [1]. In addition, within the family of the so-called ultra-high-temperature ceramics (UHTCs), ZrB2 has an important position due to its low specific mass ( 6.09 g/cm3), which makes it attractive for aerospace applications [2]. In fact, the need for future systems for a new global space access mission has led to a growing demand for high-temperature structural materials with  u Fax: +39-0546-46381, E-mail: fmonte@istec.cnr.it  improved thermo-physical capabilities in terms of strength, dimensional stability, and resistance to oxidation/ablation or to thermal shock. UHTCs have shown some potential to perform well in the environment for applications like airframe leading edges on sharp-bodied re-entry space vehicles. composites, M = Zr and Hf [3 - 5]. State-of-the-art UHTCs are represented by MB2 - SiC-based ZrB2 -based ceramics were directly fabricated from commercially available powders, or synthesized by means of reaction-based methods [6, 7]. The intentional incorporation of additives has been reported to assist beneficially the poor sinterability of ZrB2 [8, 9], but often sacrifices hightemperature strength and oxidation resistance. The integration of a SiC particulate within ZrB2 matrices provided evident profits for the oxidation and abla tion resistance [2, 3, 10 - 12], being favourable for the retention of high-temperature stability as well [2, 9, 13]. It’s the author’s considered opinion that the appropriateness of sintering temperatures for the ZrB2 - SiC system above studies of undoped MB2 - SiC, M = Zr or Hf , were primarily C has often not been considered properly. A number of addressed to achieving full density, but the influence of the grain coarsening (derived from excessive heating) adversely eroded the expected improvements in performance [4, 13]. The present study describes some interesting advances in the sinterability and mechanical performances of an additivefree ZrB2 + 10 vol. % ultra-fine α-SiC developed by hot pressing at 1900 C and 20-min dwell time, which represent more viable conditions than others reported for the same undoped system [14].  2000        2  2.1  Experimental procedure  Processing of the powder mixture  A ZrB2 + 10 vol. % α-SiC mixture was prepared from commercially available powders. The as-received powders were loaded in due proportions into a polyethylene bottle, milled for 24 h in absolute ethanol using zirconia balls, dried in a rotating evaporator under a continuous stream of nitrogen, and finally sieved ( size). The ultra-fine α-SiC fraction was dispersed ultrasonically in absolute ethanol before batching it into the final mixture. The mixture will be referred to as ZS in the rest of the paper. Table 1 lists some characteristics of the raw powders.  250-µm mesh  \\x0c', '330  Applied Physics A - Materials Science & Processing  Phase (symmetry)  Producer  Grade  Size  ZrB2 (hexagonal) α-SiC (hexagonal 6 H)  H.C. Starck GmbH, Germany H.C. Starck GmbH, Germany  B  UF25  ∗ µm: Fisher size (APS)  D90  (µm)  4-6  ∗  0.8  Main impurities (wt. %)  TABLE 1  Characteristics of the starting raw powders (producers ’ datasheets): granulometric size distribution (D90 ), specific surface area (s.s.a.)  O 2, Hf 0.2  s.s.a. (m2 /g)  −  23-26  O 2.5  The ZS powder mixture was hot pressed at low vacuum (0.5 mbar) using an inductively heated graphite mould lined with a BN-sprayed graphitized 0.75-mm-thick foil. The highest sintering temperature (measured with a pyrometer focused on the graphite mould) and the dwell time were 1900 C and respectively. The external pressures during heating and over the 20-min isothermal stage were 40 and 50 MPa, respectively. The linear shrinkage of the cold-compacted powder pellet was recorded by measuring the displacement of rams.  20 min,     2.2  Material’s characterization and oxidation test  Cu K α  The bulk density of the as-sintered billet was measured applying Archimedes’ method and using water as the immersing liquid. The crystalline phases were identified by means of an X-ray diffractometer (XRD, Ni-filtered radiation, model D500, Siemens, Germany). The microstructure of the hot-pressed material was investigated using a scanning electron microscope (SEM, model S360, Leica, Cambridge, UK) equipped with an energy-dispersive X-ray microanalyser (EDX, model INCA Energy 300, Oxford Instruments, UK). Fractured and polished surfaces of the dense material were imaged via SEM using secondary electrons (SEs). Polished cross sections were prepared applying the widely accepted metallographic method. Young’s modulus ( E ) and Poisson’s ratio ( ν) were measured on a 28.0 × 8.0 × 0.8 mm3 plate using the resonance frequency method. Flexural strength (σ ) in a four-point configuration was tested in ambient air at room temperature (five specimens tested) and on 25.0 × 2.5 × 2.0 mm3 chamfered bars using 20 mm and C (three specimens tested) 10 mm as outer and inner spans, respectively, and a crosshead speed of 0.5 mm/min. Fracture toughness (K Ic ) was measured through the chevron-notched beam method (three tested) using 25.0 × 2.5 × 2.0 mm3 bars on the specimens same jig used for the flexural strength (cross-head speed 0.05 mm/min). The thermal expansion behaviour was evaluated up to 1300 C/min heating rate) in a stream of argon using a dilatometer (model DIL 402E, NETZSCH Ger ätebau GmbH, Germany). The response to oxidation was verified by conducting an isothermal treatment at 1450 C for 20 h in flowing dry air C/min of heating rate, and free cooling, using a thermo-gravimetric analyser (model STA449 Jupiter, −3 mg of accuNETZSCH Gerätebau GmbH, Germany, racy) equipped with a vertically heated Al2O3 chamber. The specimen with dimensions 2.5 × 2.0 × 12.0 mm3 was rinsed in boiling acetone, dried at 80 C overnight, and then placed  (30 cm3/min), 30  1500  C (5  10                    upon zirconia supports with minimal contact area. The dimensional changes of the specimen were measured after the hot exposure using an optical microscope. The microstructural damage caused by the oxidation attack was evaluated via SEM-EDX examining a polished cross section. The cross section was polished to 0.25 µm using non-aqueous lubricants in order to preserve any boron-containing glass species. The results were utilized for an empirical estimation of the resistance to oxidation.  3  3.1  Results  Densification behaviour and microstructure     The bulk density of the ZS composite was 5.81 g/cm3 , which corresponds to 100% of the theoretical density estimated with the rule-of-mixture. SEM observations of the polished section confirm ed that residual porosity is not appreciable. Table 2 shows some microstructural attributes of the ZS material. Figure 1 plots the densificatio n curve during hot pressing. The onset of some measurable shrinkage of the green body during hot pressing was associated with the temperature TON (Table 2). A marked increase in the densification rate occurred only for temperatures higher than 1500 C. The densification curve of the same ZrB2 pure powder was added to Fig. 1 for comparison. The XRD analysis of ZS material, besides ZrB2 and SiC, did not detect any extra crystalline phases. The fracture surface examined by SEM typically shows a grained structure, with regularly faceted ZrB2 grains and the SiC particles dispersed within the diboride skeleton (Fig. 2). The grain-size distribution of the ZrB2 matrix is narrow (1 - 4 µm), and the average size does not exceed 3 µm. The fracture appears to propagate predominantly along an intergranular boundary interfaces are apparently free of secondary foreign phases. The SiC particulate is mainly distributed intergranularly, although a limited fraction remains located intragranularly. In addition, sev path. The ZrB2 /ZrB2  TON     C)  (  THP     C)  (  t  (min)  Density  Crystalline phases  d (g/cm3 )  ∗  rd  Main  Secondary  1260  1900  20  5.81  100%  ZrB2 , SiC  None  ∗  estimated by SEM on polished surface  mgs (µm)  = 3  TABLE 2  Onset and hot-pressing temperatures (TON , THP ), dwell time (t ), bulk (d ) and relative density (rd), crystalline phases, and ZrB2 mean grain size (mgs)  \\x0c', 'MON T EV E RD E Beneficial effects of an ultrafine  α-SiC incorporation on the sinterability mechanical properties of ZrB2  331  FIGURE 1  Relative density (rd) vs. time (t ) of the ZS material during hot pressing. The densification curve from pure ZrB 2 was added for comparison [9]. The dotted portions of the densification curves cover the isothermal stage  FIGURE 2  material  Micrographs by SEs-SEM from fracture  surfaces of  the ZS  eral original contact points between ZrB2 grains and SiC particles are observable. From the investigation by SEM of the polished section, neither residual porosity nor reaction by-products of the hot pressing were found in appreciable amounts.  3.2  Thermo-mechanical properties and resistance to oxidation     Table 3 summarizes the experimental data for the some equivalent properties of a ZrB2+10 vol. % SiC material measured thermo-mechanical properties. For comparison, hot pressed at 1900 C and 45-min dwell time [5], or other ZrB2 compounds, are shown. the ZS composite, 507 ± The Young’s modulus ( E ) of agreed with the value expected from the rule-of mixture E<ZS> = V1 E<ZrB2> + V2 E<SiC> , where V1 and V2  are the volume fractions of ZrB2 and SiC, respectively. Values  4 GPa,  for E<ZrB2> and E<SiC> were 530 GPa [15] and 440 GPa [16],  FIGURE 3  Load ( L ) vs. displacement (d ) curve during flexural strength tests of the ZS material at 1500 C. Linearity of the terminal part of the curve is apparent        1300  respectively. Residual porosity and low-density secondary phases are known to affect this property unfavourably. In the present case, the measured E value proved to be essentially controlled by the elastic constants of the pure compounds composing the hot-pressed ZS material. In addition, Poisson’s ratio (ν) did not shift appreciably from that of the monolithic ZrB2 (i.e. 0.11) [17]. Linear behaviour describes the thermal expansion up to C. The linear thermal expansion coefficients (LTECs) −6 /K fre(Table 3) range within the interval 7.0 - quently mentioned for the ZrB2 - SiC system [9, 13]. The fracture toughness (K Ic ) and the flexural strength ( σ ) received an advantageous contribution from the incorporated SiC particulate, and demonstrated an interesting combination of values for this category of materials. The optimization of the fabrication method, the controlled sintering temperature for instance, caused the refinement of the ZrB2 matrix (about 3-µm average grain size), which, in turn, effected a significant improvement of th e room-temperature strength (Table 3).  7.5 × 10     at 1500  At 1500  C, a reduction in strength of about 64% of the initial value occurred. However, the load-displacement curve C did not deviate from linearity appreciably (Fig. 3). The reasons for such a decrease will be commented on in the discussion section. The resistance to oxidation was investigated over the 20-h thermo-gravimetric (TG) experiment C. Figure 4 shows the graphical trend of the mass gain vs. time: the final  at 1450        Material  p  E  ν  (vol%)  (GPa)  # K Ic  √  (MPa  m)  ZS ZrB2 -10 SiC [5] ZrB2 [9] ZrB2 [25] # mean ±1 st. dev.;  ∧  None 6.8 13 None  507 ± 4 346 ± 4 450 550  ∗  0.12  − − −  4.8 ± 0.3 4.1 ± 0.3 2.35 ± 0.15  −  unpublished datum;  ∗  three-point bending  (25-1000)  7.35  −  7.54  −  LTECs (10 C)  −6/      C  (25-1300)  7.57  −  7.51  −     C     C  25  835 ± 35 713 ± 48 350 ± 30 300 ± 20  ∗  # σ  (MPa)     1400  C  − −  220 ± 10 200 ± 50  ∧ ∗     C  1500  300 ± 35  − − −  TABLE 3  Porosity ( p), Young’s modulus ( E ), Poisson’s ratio ( ν), fracture toughness (K Ic ), linear thermal expansion coefficients (LTECs), and four-point flexural strength ( σ ) in air at various temperatures  \\x0c', '332  Applied Physics A - Materials Science & Processing     left FIGURE 4  (KD , right-hand Y axis) vs.  Weight change (w, hand Y axis) and dummy parabolic constant time (t ) over the oxidation test at 1450 C for 20 h. The shows the heating ramp up to 1450 C (30 C/min heating rate). The indicate discontinuities of log(KD ) curve (see the text)  rows  inset        ar the        C and 1500  C for 20 h had 5 and 10.9 mg/cm2  mass gain of the oxidized sample was 4.80 ± 0.02 mg/cm2 . An offset equal to 0.65 ± 0.02 mg/cm2 , which accounted for from the TG data. For comparison, ZrB2+20 vol. % SiC oxithe oxidation preceding the fixed exposure, was subtracted dized at 1400 of weight gain, respectively [12]. At first sight, the TG data in Fig. 4 state that the oxidation mechanisms involved at these conditions were partly protective. Data analysis in fact was carried out using a multiple-law model proposed by K.G. Nickel [18]. A dummy parabolic rate constant K D was calculated over the 20-h exposure for each datum point assuming that K D = w/ t (w is the weight change, t the time), and then plotted in Fig. 4. Such an exercise readily points out that a passivating parabolic kinetic (i.e. slope of the log(K D ) data curve equal to zero) had never been established. Instead, the paralinear law  √  √  w = K PAR  t + K LIN t  FIGURE 5  Micrograph by SEs-SEM (polished cross section) of the specimen oxidized at 1450 C for 20 h: the outermost glassy coating (a), the poorly compact oxide scale (b), and the unoxidized bulk (c)     (K PAR and K LIN parameters) was in very good agreement ( R2 > 0.995) with the mass-gain data up to about 500 min of exposure, with the best-fitting K LIN parameter negative (Fig. 4). Afterwards, a linear behaviour adequately fitted the weight change for the remaining exposure time. The pattern of the log(K D ) data curve, along with some discontinuities (see arrows in Fig. 4), was connected to a progressive decay of the external oxide scale protectiveness, until a non-passivating linear kinetic took place. Specific discussion is presented later in the paper. The examination by SEM-EDX of the oxidized sample agreed with the remarks formerly made. The polished cross section highlighted a multi-layered arrangement of differently oxidized regions (Fig. 5). An outermost glassy layer thick) covers a ZrO2 sub-scale (about 200-µm thick). Isolated ZrO2 crystals are present inside the external coating. The adherence of the glassy layer to the oxide sub-scale appears con (5-µm  tinuous, though cracks and a small amount of detachment at the interface of the two scales were observed (Fig. 6). Microcracking was reasonably assumed to have originated during cooling because of the thermal expansion mismatch between the silica glass and zirconia. Some localized detachment due to the polishing cannot be excluded. The XRD analysis of the oxidized surface confirmed the formation of highly textured 00l monoclinic zirconia crystals. The EDX analysis of the external glassy coating, Si and O apart, did not detect Zr or C. As far as the presence of B is concerned, a limited amount incorporated in the external glass layer was found (Fig. 6). Poorly dense regions compromise the compactness of the ZrO2 sub-scale below the borosilicate glassy coating, chiefly in the vicinity of the external surface. Such an intermediate layer extends up to the unoxidized bulk. Multiple sites within  \\x0c', 'MON T EV E RD E Beneficial effects of an ultrafine  α-SiC incorporation on the sinterability mechanical properties of ZrB2  333  FIGURE 6  Left-hand side: micrographs by SEs-SEM (polished cross section) of the specimen oxidized at 1450 external oxide scale); right-hand side: EDX spectrum (electron-beam energy 3 keV) of the external glass layer     C for 20 h (a magnified view from the  FIGURE 7  Left-hand side: micrograph by SEs-SEM (polished cross section) from the zirconia sub-scale of the specimen oxidized at 1450 C for 20 h: the arrows mark some C-containing inclusions; right-hand side: EDX spectra (electron-beam energy 3 keV) from these inclusions, and from SiC for comparing the compositional transition     this oxide sub-scale, which were formerly occupied by the SiC particulate, appear in the form of inclusions filled basically with carbon (Fig. 7). This statement is further supported by the evidence of the size, shape, and distribution of these features, which are referable to the original SiC particulate  in the virgin material (Fig. 8). In contrast to other studies of the (undoped) ZrB2 - SiC system [11, 12], the disappearance of the SiC particles by active oxidation unexpectedly resulted in the formation of a condensed carbon-based by-product instead of their complete volatilization. The dimensional measurements of the oxidized specimen indicated an increase in the initial size of about 50 µm. This result confirmed that, even t hough mass loss occurred to some extent, the whole oxidation process did not involve a bulk recession.  4  4.1  Discussion  Densification and microstructural texture  The densification behaviour of the ZrB2 powder was greatly improved by adding the ultra-fine α-SiC particulate. ZrB2 powder on its own, processed in the same way [9], exhibited a sinterability significantly lower than that of the composition herein tested: in the former case, the final relative density was 87% of the theoretical value (Fig. 1). Chamberlain and co-workers reported a value of relative density of  93.2% for ZrB2+10 vol. % α-SiC (0.7-µm mean size for SiC),  unintentionally doped with 2 vol. % WC during milling and hot pressed at 1900 C for 45 min and 32 MPa of applied pressure [5]. Another research group, which had worked exten    FIGURE 8  Micrograph by SEs-SEM (polished cross section) of the specimen oxidized at 1450 C for 20 h: the inner unoxidized bulk     \\x0c', '334  Applied Physics A - Materials Science & Processing              sively on additive-free ZrB2 - x SiC compositions, 10 < x < 30 and 5-µm mean size for SiC, addressed the most important priority to the full densification, accepting hot-pressing conditions of temperatures up to 2150 C for 2 h [19]. In this sense, the incorporation of the ultra-fine α-SiC particulate fraction herein presented, instead of more widely used coarser powders, enabled full density to be achieved by applying less severe sintering conditions. Regarding the shrinkage rate from 1260 C (i.e. TON ) up to about 1500 C (Fig. 1), the densification of the ZS material was supposed to be controlled by some rearrangement of the ultra-fine SiC powder fraction. Aided by the applied pressure, a slow flux of the SiC particulates to the pores at the edges of the ZrB2 grain faces presumably took place until the powder compact reached a fractional volume of porosity equivalent to that of other powder mixtures similarly processed [20]. Above 1500 C, densification started speeding up. In comparison with the pure ZrB2 , densification of the ZS powder mixture proceeded more efficien tly, making clear the relevant role of α-SiC in improving the sinterability of ZrB2 . In-air storage and processing of a metal diboride powder like ZrB2 are known to contaminate its particle surface with oxygen. Such contamination by oxygen of powders favours, during the hot consolidation, mechanisms like evaporation-condensation a nd grain coarsening, which worsen the maximum attainable density. In the present case, the attainment of a refined fully dense ZrB2 matrix clearly verified the beneficial role SiC played in limiting deleterious mechanisms for densification, like evaporation-condensation (Fig. 9). The way in which the addition of an ultra-fine SiC particulate assisted the full densification of ZrB2 so favourably has not been completely clarified. Some authors have asserted the cleaning of the ZrB2 surface powder particles from oxygen as the primary step for obtaining a highly dense body [9, 21]. It therefore seems plausible, also for the current case, to look for a depletion in the oxygen impurity of ZrB2 through chemical interactions with SiC. Specifically, SiO2 and B2O3 were assumed as the main oxygen carriers upon the surfaces of SiC and ZrB2 , respectively. In addition, the absence of solid secondary phases in the hot-pressed ZS led us to assume the evolution of only gaseous by-products between SiC and the oxygen-bearing compounds. The reaction  the ZrB2  SiC + B2O3 + SiO2 = SiO(g) + CO(g) + BxO y (g) . . .  (1)  is (thermo-chemically) favoured at the processing conditions. The multi-phase equilibrium in Fig. 10 tentatively depicts the consumption of the oxygen-containing contaminants at temperatures below those associated with the beginning of a significant shrinkage of the cold-compacted powder mixture (Table 2). Such a circumstance permits the volatile byproducts to evolve outside through interconnected open channels of the ‘green’ compact. The removal of the oxygencarrier species from the ZrB2 surface particles had the effect of increasing the boron activity through the generation of Zr vacancies [21]. A higher metal-vacancy concentration restores a favourable requisite for lattice diffusion, and thus for densification. The application of increasing temperatures meets the conditions for boundary diffusion, and assists densification efficiently, the boron activity being depressed no further. Some intragranular SiC particles inside ZrB2 grains provided evidence of the occurrence of a significant mass transfer among the diboride particles, which in effect brought the complete removal of residual porosity to a successful conclusion. Compared with a ZrB2 mixture containing coarser α-SiC particles [5], the present system likewise processed at 1900 C showed a more pronounced tendency towards sintering (Table 3). In fact, the ultra-fineness of stimulated a higher reactivity in the system studied, helping reaction (1) to advance more effectively.  10 vol. %  α-SiC     4.2  Thermo-mechanical properties     ZrB2+10 vol. % SiC system hot  The comparative list of the experimental data in Table 3 emphasizes the performance merits gained by the addition of a SiC particulate to ZrB2 . To be specific, the assistance of an ultra-fine ceramic component had succeeded in the full densification of the undoped pressed at (only) 1900 C. It implies that, the complete removal of porosity apart, grain-size refinement and control of secondary phases, which are well-known key requirements for the improvement of the material’s properties, were both achieved successfully. This undoubtedly was positively reflected by enhanced mechanical properties like strength, fracture toughness, and Young’s modulus in the material herein developed.  a  b  FIGURE 9  Micrographs by SEs-SEM from the fracture surface of the ZS material (a) and the pure ZrB2 (b), both hot pressed at 1900     C  \\x0c', 'MON T EV E RD E Beneficial effects of an ultrafine  α-SiC incorporation on the sinterability mechanical properties of ZrB2  335  FIGURE 10  Isobaric p = 0.5 mbar multi-phase equilibrium calculated at [22]; the starting composition (mol) 1 SiC, 1.1 SiO2 , and 0.9 B2 O3 was chosen for convenience     √  of another comparable ZrB2+10 vol. % α-SiC [5] relate how Data referred to the microstructure and some properties adverse is the densification of t he mentioned system at a temperature of about 1900 C. Young’s modulus exceeding 500 GPa and fracture toughness of about 5 MPa m are valuable results. On one hand, these two properties are not adversely affected by the residual porosity. In addition, fracture toughness increased (Table 3) because of the presence of sites (i.e. SiC particles) in proximity of which the crack, when propagating, might experience energy-dissipating events (i.e. crack deflection /bridging), responsible for the toughening of this fine-grained composite. To better elucidate the toughening mechanisms, the paths of some Vickers-indentation-induced cracks were examined on polished sections via SEM. In Fig. 11, the SiC particles clearly appeared to deflect the crack and /or bridge its opening surfaces. The refinement of ZrB2 was already stated to lead to a significant increase in strength (at room temperature). In contrast to other doped ZrB2 - SiC materials [9], the load-displacement curve registered at 1500 C (in static air) obeyed a basically linear pattern (Fig. 3). The observation by the optical microscope of the specimens tested at 1500 C did not provide evidence either of residual plastic deformation or sub-critical crack growth. Actually, the fracture surfaces of the bars tested C had been modified consistently by oxidation after fracturing. This prevented the fractographic study of the critical flaws that initiated the failure. Nonetheless, microstructural changes just beneath the specimens ’ faces exposed to air C occurred. They consist of a partially compact zirconia layer, covered with a thin borosilicate glass layer. The formation of such a zirconia scale was looked at in trying to give a reason for the most significant contribution of the strength decrease at 1500 C. A more detailed insight into the response to oxidation was accomplished over the 20-h isothermal cycle at 1450 C in air.  at 1500  at 1500                             C  to 1370  and 1400  The undulating change in the oxidation rate between 1250 C (see inset in Fig. 4) disclosed how the incorporated SiC enhanced the resistance to oxidation of ZrB2 alone. It is well known that ZrB2 , once oxidized, yields ZrO2 and tion SiC + 3/2 O2 (g) = SiO2 + CO(g), reacts with oxygen B2O3 . However, silicon carbide, in accordance with the reacand provides condensed silica (i.e. passive oxidation), which begins to combine with the available B2O3 to form a borosilicate protective glass. Such an oxidation product retards not only the vaporization of B2O3 (i.e. net mass loss from 1250 C) but significant further inner diffusion of oxygen as well, giving more oxidation protection than zirconia alone. Such a description agrees with the paralinear pattern previously proposed: the parabolic contribution accounts for the diffusion of oxygen through the multi-phase external oxide scale, whilst the negative linear term relies on the continuous release of gaseous by-products like the boron oxides. The examination of the reaction-scale morphology revealed not only a partially porous structure composed of zirconia grains without a glass coverage on them, but carboncontaining inclusions throughout the oxide scale up to the oxide/diboride interface (Figs. 7 and 12). These carbon-based structures were associated with the active oxidation of SiC. Carbon apart, the detection of Si and O in such inclusions (Fig. 7), which in turn exhibited the size, shape, and distribution of the SiC particulate in the unoxidized material (Fig. 8), led us to consider the following thermodynamic equilibrium  SiC + SiO2  ⇔  2SiO(g) + C . . . . . .  (2)  as the principal transition mechanism by which SiC experienced the active decomposition. In the remarkable research work of Gulbransen and Jansson on the oxidation of SiC, two other thermodynamic equilibria  \\x0c', '336  Applied Physics A - Materials Science & Processing  FIGURE 11 Micrographs by SEs-SEM: (a) the path of a Vickersindentation-induced crack on a polished surface; (b) the magnified view of the feature indicated by a circle in (a). The arrow indicates the crackpropagation direction  2SiC + SiO2 SiC + 2SiO2  ⇔ ⇔  3Si + 2CO(g) 3SiO(g) + CO(g)  (3)  (4)  were investigated as potential transition mechanisms [23]. Conditions of high temperature and low oxygen partial pressure inside the oxide scale were met for the active oxidation of SiC [24]. Once SiO(g) diffuses outward and encounters a higher oxygen partial pressure, it would further convert into the condensed silica phase. Other authors reported the loss of SiC in ZrB2 - SiC materials through active oxidation with the release of only gaseous by-products [10 - 12]. The reason why a similar active oxidation of SiC evolved in agreement with equilibrium (2) is unclear. The presence of different SiC polytypes, i.e. ultra-fine α, instead of coarser β, was proposed as one possible factor promoting the decomposition in agreement with equilibrium (2). Likewise, the absence of both residual porosity and secondary grain boundary phases in the ZS material, differently from those aforemen a  b  tioned [10 - 12], constitute favourable conditions which, together with the formation of a continuous external borosilicate glass layer, limited substantially the inner diffusion of oxygen, and therefore its partial pressure PO2 inside the bulk. The PO2 parameter is recognized as the primary factor controlling the active-passive oxidation transition of SiC [23, 24]. Evidently, the active oxidation of SiC had undesired consequences on the resistance to oxidation of the studied material. The thin outermost glassy scale, along with the creation of pores, facilitated the active oxidation continuing to a greater depth, the transport of oxygen being allowed through the less protective zirconia scale. Zirconia alone is known to protect partially against oxidation at elevated temperatures, due to its anion-deficit structure which permits transport of oxygen via lattice vacancies [10]. The persistent departure of the mass-change data from a parabolic pattern (Fig. 4) implies that, at these oxidizing conditions, effective protection had never been established completely. The discontinuities of the log(K D ) data curve (emphasized in Fig. 4) were supposed to derive from the cyclic build-up of the inner pressure (connected to evolved volatile species) that partially disrupts the cohesiveness of the oxide scale formed. The subsequent opening of new accesses for the inward diffusion of oxygen, along with the thinness of the outermost glassy layer, rendered the external oxide scale only semi-protective for the unoxidized parts of the material. The linear trend of the weight-change data above 500 min of exposure accounts for the transition into a linear non-protective oxidation regime. The depletion of SiC has such an important impact that it restricts the applicability of this material for long-term thermally severe applications.  5  Conclusions     The present work outlines some interesting advances in the densification and thermo-mechanical properties of fully dense ZrB2 , hot pressed at 1900 C with the addition of 10 vol. % ultra-fine α-SiC particles. The incorporated α-SiC particulate played a key role in enhancing the sinterability of ZrB2 , expressly assisting the removal from the diboride surface particles of the oxygen impurity, which severely limits the maximum attainable density. The microstructure consisted of fine-faceted ZrB2 grains, and SiC particles dispersed evenly throughout the diboride framework. No extra secondary phases were found. The SiC  (a) the crossways dotted line at  FIGURE 12 Micrographs by SEs-SEM from the (polished) cross section of the specimen oxidized at 1450 C for 20 h: the oxide scale/virgin bulk interface; (b) a partially active-oxidized SiC particle enclosed by carbon by-product     \\x0c', 'MON T EV E RD E Beneficial effects of an ultrafine  α-SiC incorporation on the sinterability mechanical properties of ZrB2  337  phase inhibited excessive coarsening of ZrB2  3 µm).  (average size  (cid:8)  s     m  √  4.8 ± 0.2 MPa  300 ± 35 MPa at 25 and 1500 toughness, 0.12 Poisson  If compared with monolithic ZrB2 , the mechanical properties of the composite being tested were of 507 ± 4 GPa Young’s modulus, relevant significance: and 835 ± 35 and fracture ratio, C in-air flexural strengths, respectively. The ZrB2+10 vol. % α-SiC material did not prove to endure oxidation conditions acceptably at 1450 C for 20 h in (flowing) air. Even though lim ited mass gain was measured and no dimensional recession was observed, effective protection against the oxidation attack was never established. In particular, active oxidation of SiC and its conversion into carbon-based inclusions were ascertained. This result makes the structural capability of this composite not properly suited for long-term use at very high temperatures in oxidizing environments.     ACKNOWLEDGEMENTS  The author wishes to acknowledge the helpful assistance of his colleagues, including D. Dalle Fabbriche (hot pressing), C. Melandri (mechanical characterization), and A. Balbo (thermal expansion and oxidation tests).  REFERENCES  1 C. Mroz: Am. Ceram. Soc. Bull. 73, 141 (1994) 2 K. Upadhya, J.M. Yang, W.P. Hoffman: Am. Ceram. Soc. Bull. 58, 51 (1997) 3 M.M. Opeka, (2004) 4 M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. 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A 4, 2661 (1986) 16 MCIC-HB-07 - Vol II Battelle Institute: Engineering Property Data on Selected Ceramics, Carbides (Battelle Laboratories, Columbus, OH 1987) pp. 1-19 17 R.A. Cutler: Engineered Materials Handbook, Vol. 4, ed. by S.J. Schneider (ASM International, Materials Park 1991) p. 787 18 K.G. Nickel: Modelling, ed. by K.G. Nickel (Kluwer, Dordrecht 1994) p. 59 19 E.V. Clougherty, L. Kaufman, D. Kalish: US Patent No. 3 775 137 (1973) 20 A. Bellosi, F. Monteverde, G.N. Babini:  in Engineering Ceramics ’96: Higher Reliability Through Processing, ed. by G.N. Babini, M. Haviar,  in Corrosion of Advanced Ceramics/Measurements and  Peters: Tech. Rep. Laboratory, Wright P. Sajgalik (Kluwer, Dordrecht 1997) p. 197 21 D.N. Øvrebø, F.L. Riley: Conf. Exhi b. 6th ECerS, Ext. Abstr. Vol. 2, Br. Ceram. Proc. No. 60, 19 (1999) 22 A. Roine: HSC Chemistry 5.11, Outokumpu Research Oy, Pori, Finland (2002) 23 E.A. Gulbransen, S.A. Jansson: Oxid. Met. 4, 181 (1972) 24 W.L. Vaughn, H.G. Maahs: J. Am. Ceram. Soc. 73, 1540 (1990) 25 D. Kalish, E.V. Clougherty, K. Kreder: J. Am. Ceram. Soc. 52(1), 30 (1969)  \\x0c']"
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  "_id": 24,
  "PDF": "Carbon fiber reinforced hafnium carbide composite..pdf",
  "Text": "['J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5995 - 6003  Carbon fiber reinforced hafnium carbide composite  ULTRA-HIGH TEMPERATURE CERAMICS  A . S A Y I R NASA Glenn Research Center/Case Western Reserve University, Cleveland, OH 44135, USA E-mail: Ali.Sayir@grc.nasa.gov  Hafnium carbide is proposed as a structural material for aerospace applications at ultra high temperatures. The chemical vapor deposition technique was used as a method to produce monolithic hafnium carbide (HfC) and tantalum carbide (TaC). The microstructure of HfC and TaC were studied using analytical techniques. The addition of tantalum carbide (TaC) in the HfC matrix was studied to improve the microstructure. The microstructure of HfC, TaC and co-deposited hafnium carbide-tantalum carbide (HfC/TaC) were comparable and consisted of large columnar grains. Two major problems associated with HfC, TaC, and HfC/TaC as a monolithic are lack of damage tolerance (toughness) and insufficient strength at very high temperatures. A carbon fiber reinforced HfC matrix composite has been developed to promote graceful failure using a pyrolytic graphite interface between the reinforcement and the matrix. The advantages of using carbon fiber reinforcement with a pyrolytic graphite interface are reflected in superior strain capability reaching up to 2%. The tensile strength of the composite was 26 MPa and needs further improvement. Heat treatment of the composite showed that HfC did not undergo any phase transformations and that the phases comprising composite were are thermochemically compatible. C(cid:1) 2004 Kluwer Academic Publishers  1.  Introduction              Expandable and reusable space vehicles, next generation rocket engines and hypersonic spacecraft need materials and structural components capable of operating at temperatures in excess of 1600 C. The ultimate payoff is expected to come when materials are developed that can perform without cooling at gas temperatures exceeding 2200 C. Temperatures above 1600 C, and if possible exceeding 2200 C, will be described as the ultra high temperature region to differentiate the unique thermomechanical and thermochemical demands of aerospace applications. Materials for rocket combustion chambers, thrusters, and nozzles must meet several requirements simultaneously, such as high melting temperature, minimum strength, and environmental resistance (i.e., oxidation resistance). The selection of potentially suitable structural materials for use at ultra high temperatures in air breathing engines was identified in reports by the Air Force Wright Aeronautical Laboratories [1-6]. Borides, carbides, boride-graphite composites, carbide-graphite composites, pyrolytic and bulk graphite, coated refractory metals/alloys, oxide-metal composites, oxidation resistant refractory metal alloys, oxide-metal composites, and iridium-coated graphite were considered for ultra high temperature applications (see Table I). Boride composites were developed at Air Force Laboratory for leading edge and nose cap applications which require moderately high temperatures. Several other researchers proposed monolithic ZrB2 [7-9] and HfB2  [7, 8, 10], as well as and composites of these boride materials with other ceramics, including TiC/ZrB2 , TiC/HfB2 , SiB6 /ZrB2 , [8], and ZrB2 /SiC or HfB2 /SiC. A major reason for proposing these materials, however, was the availability of easy fabrication through hot pressing, resulting in high relative density and a concomitant high strength. The oxidation resistance of these materials was limited [11-15]. A renewed effort to fabricate HfB2 -SiC and ZrB2 -SiC composites is underway to tackle the challenges related to oxidation that exist for ultra high temperature region [16-18]. Still, the development of processing technology to produce near-net shape components is a formidable task. A second group of important materials for aerospace applications is based on monolithic refractory metal and refractory metal based composites [19]. These materials are of interest for their high strength and ductility at elevated temperatures. Some of these materials should be capable of carrying significant loads at ultra high temperatures, but they have the disadvantage of very low specific strength. Engine weight and inertial force considerations often put a premium on utilizing materials with low density. Hence, the mechanical requirements, such as strength, stiffness, and creep resistance have to be normalized with respect to density. In rotating or reciprocating machines, and especially in structures to be airor space borne, these density-normalized properties become the criteria of interest. From the point-of-view of specific strength requirements the metals listed in Table I (W, Mo, Ir,  0022-2461  C(cid:1) 2004 Kluwer Academic Publishers  5995  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E  I Classification of materials using single criteria; melting  point [1-10, 20-22]  Type of material  Carbon  Metals  Intermetallics  Light Ceramics  Refractory metal  Ceramics  Prime materials     Tm > 3000  C  Marginal materials Tm (cid:3) 3000  C     Diamond, graphite  W, Mo, Ir, Os, Ta  ReW, Re3W, Re2 Ti5 , SiC, B3 Si HfC, TaC, NbC, ZrC, WC, W2C, VC,MoC, Ta2C, TiC, HfN, Mo2C, ZrN, TiN, WB, TaN, HfB, HfB2 , TiB2 , Nb3B4 , WB2 , YB2 , TaB2 , ZrB2 ,NbB2 , ZrB, W2B, HfO2 , UO2 , ThO2 ThZrO4 , MgO, ZrO2 Er2O3 , ZrO2 , SrZrO3  Os and Ta) are not satisfactory. These metals, W, Mo, Os and Ta, also have very poor oxidation resistance. Iridium and compatible alloys perform admirably for rocket throat and nozzle applications [20-22], but high density, high cost of raw materials, and laborious machining operations makes iridium and its alloys less attractive. A third group of candidate materials is based on high temperature carbides and nitrides. In ultra high environments such as that of rocket engines, structural components must be capable of withstanding shock and high structural loads in highly corrosive environments. A material therefore should not undergo any significant surface degradation in high temperature oxidizing environments. This requirement eliminates from consideration all materials that react with oxygen to form volatile products. The volatility can be quantified in terms of vapor pressure and oxidation-induced surface recession rates as a function of temperature. Shaw et al. [23], Wicks et al. [24] and Shick et al. [25] have shown that the carbides tend to have lower pressures than borides and nitrides. Using the volatility data from Shaw et al. [23], one can conclude that HfC is one of the candidate materials for use in ultra high temperature environments. Another primary requirement of a candidate material is low diffusion coefficients at high temperatures. The material should be stable with respect to chemical interactions with the oxidation product, as well as any secondary phases present, over long periods at elevated temperatures. These objectives can be met by selecting materials that possess large negative free energies of formation. Hafnium carbide fulfills these thermodynamic requirements and it is the choice of material for this investigation. Another justification for the selection of HfC is derived from the mode of oxidation of HfC to HfO2 . During oxidation of hafnium carbide, a very distinctive and heterogeneous structure forms that has important implications in potential applications [26-30]. The structure contains three distinct layers: (a) a residual carbide layer with dissolved oxygen in the lattice, (b) a dense oxide interlayer containing carbon (HfCx O y ), and (c) a porous outer layer of hafnium oxide. Bargeron et al. [31] has shown that the inter( Di,eff = 1.1 × 10 layer HfCx O y has remarkably low diffusion coefficient −7 cm2 /s at 2060 C) and is an oxygen diffusion barrier.     5996  Although chemically stable at high temperatures, HfC would require fiber reinforcement to attain strainto-failure capability (increased toughness) suitable in load-bearing applications. High strength carbon fibers are currently the only suitable reinforcement for the composites used at ultra high temperatures. Carbon fibers can readily be formed into a construct or perform of desired configuration by winding, weaving, knitting, braiding, or wrapping over a suitably formed mandrel. In addition, carbon fibers are relatively inexpensive. A composite of hafnium carbide as the matrix and carbon fibers as the reinforcement emerges as the candidate material having the highest upper use temperature. The aim of this work was to produce carbon fiber-reinforced hafnium carbide (HfC) composites and study the microstructure—mechanical property relationship. The goal is to produce a strong and tough composite that will withstand an operating temperature of 2200 C but for short time applications (minutes rather than hours). The HfC matrix was produced via chemical vapor deposition (CVD). This method provides a high degree of control of deposition rates and is capable of producing complex shaped components with superior uniformity.     2. Experimental  Pyrolytic graphite, HfC, TaC, and HfC/TaC were deposited by CVD on different substrates by Advanced Ceramics Corporation (ACC)1 . The CVD process was carried out using an industrial-scale furnace which consisted of a graphite susceptor, insulation jacket, and induction coil which were contained within a waterjacketed steel vacuum shell that could be evacuated by a vacuum pumping system. It is significant to point out that geometrical parameters play a dominant role in CVD processes and thus, the materials produced in this study would be most comparable to other materials produced by industrial-scale operations. The CVD processing conditions of carbide deposition were proprietary information of ACC and the study of processing conditions was beyond the scope of the present investigation. Graphite and two dimensional (2D) square-weave carbon fabric (trade name WCA graphite cloth produced with Rayon fibers) were supplied by POCO Graphite Inc., and Morgan Specialty Graphite, respectively. The graphite cloths consisted 10 strands per centimeter and had tensile strength of 150 MPa. Pyrolytic graphite was deposited as an interfacial layer between the carbide (HfC, TaC, or HfC/TaC). It was possible to produce Cfiber /HfCmatrix composites with complex shape. Wrapping woven carbon fiber cloth over a suitably formed support or mandrel produced the complex shape. Two-dimensional (2D) woven carbon fabrics were coated with 5 to 120 µm thick layers of pyrolythic graphite (PG) on various graphite mandrels. There were several reasons for the deposition of PG on 2D woven carbon fabric. This process rigidized the fabric, thereby  1 1197 Lakewood, Ohio (Now GE Advanced Ceramics, Strongsville,  Ohio).  \\x0c', 'producing components that could be handled conveniently without bending of the woven carbon-fiber fabric and which retained the shape of the mandrel. The rigidization of fiber fabric through PG deposition had a profound effect on the mechanical properties of the composite, as discussed later. The PG interface layer (5 µm) protected fibers from the reaction with the HfC matrix, thus minimizing fiber strength degradation during processing. PG interface resulted in porous HfC or TaC layer between dense HfC or TaC matrix. Additionally, stereological characterization suggests that the composite was approximately 60% dense. The processing of complex-shaped structures, however, produced unattractive tensile specimens, which had curvatures. Ribbons of approximately 1×10 mm cross-sections and 20 cm in lengths were machined from these composites with complex shape and curved sections. Tensile strength data were collected using these specimens, although they had an undesirable and inhomogeneous stress state. The assignment of the stress values was estimated from cross-section measurements and did not take into account the torsion and bending forces during the test. The primary intent here was to have a good engineering estimate of the strain capacity of the composite. Accurate axial strain measurements were used to assess the performance of the composite, specifically the strain-to-failure capability. In contrast to the stress measurements, the strain measurement in axial direction was very accurate. For the tensile tests, cold grips were used with a total gauge length of approximately 40 mm. All samples fractured within the 40 mm gauge length. The strain rates were calculated from the total gauge length and crosshead speed of the test frame (Model 4502, Instron Corp., Canton, MA). A computer controlled digital extensometer (Universal Dimension Meter; UDM500A from Zimmer Corporation) was used to measure the extension. The universal dimension meter (UDM) is specifically designed for precision strain measurements and uses a Mercury Xenon illuminator with suitable collimating lens that illuminates the sample from behind. Flags of silicon carbide monofilaments (SCS6/Textron Corporation) were attached to the sample in order to provide a contrasting edge. The distance between the flags was the gauge length and it was measured with a micrometer. The UDM measures edge positions without contact using a digital line scan method. Its resolution is 5000 pixels at a scan rate of 1800 Hz. The strain measurement system was ASTM83-92 Class A. The measurement error of strain was less than 2% at 25 mm gauge length and expected to be lower at longer gauge lengths. The resolution was 0.25 µm with a maximal sampling rate of 1000 Hz. X-ray diffraction analysis was carried out using a two-dimensional (2D) General Area Detector Diffraction System, GADDS, manufactured by Bruker AXS, Inc (formerly Siemens AXS). It is a true proton counter in a large area and has a fast rate of data collection. For example, the speed of data collection with an area detector can be 104 times faster than with a point detector and about 100 times faster than linear position-sensitive detector. Phase identification has been done by integrat ULTRA-HIGH TEMPERATURE CERAMICS  ing over a selected range of the Bragg angle (2\\x01) and the azimutal-angle (χ ) about the direction of the incident X-ray beam. The GADDS is also equipped with special flat graphite based rotating-optics monochromators, which produce the strongest beam intensity. The graphite monochromators cannot, however, resolve Kα1 and Kα2 lines. Hence, the system was aligned to Kα line to accomplish maximum intensity. The focal spot and critical angle are important features of the system. take off angle with a 1 × A 0.8 mm collimator at 13 10 mm focal spot size at the anode generates the maximum targeting load and brightness. The GADD system operates with a Siemens Kristalloflex 760 X-ray generator which was set at 40 KV and 40 mA and remained stable during the course of this study. Scanning electron microscopy (Joel 840, UTW Kevex EDS) was employed to characterize microstructures and to identify phases.     3. Results and discussion  Initial CVD trials were conducted to investigate the microstructures of HfC and TaC samples deposited on POCO graphite and pyrolytic graphite (PG) substrates. SEM analysis indicated that the microstructures for both POCO graphite and PG substrates were similar and remaining of the work was completed using PG substrates. Figures 1 and 2 are representative SEM micrographs of HfC and TaC samples, respectively. The SEM analysis indicated that both HfC and TaC had a gradually changing structure and had thicknesses of 120 and 30 µm, respectively. HfC contained some voids (Fig. 1a) and the microstructure was void-free as thickness of the coating increased to 30 µm, Fig. 1b. TaC did not contain any void and cracks even though the sample had to be machined and polished extensively to reveal the microstructure, Fig. 2. The CVD-deposited material tends to grow in a columnar structure. Although the strength of the individual grains may be high, the intergranular strengths can be quite low due to poor cohesion at the boundaries [32]. Average grain sizes exceeding 10-15 µm were produced without microcracks. HfC deposition on pyrolytic graphite (PG) nucleates at the grain boundaries of the PG as small and randomly oriented grains and these random nucleation sites produce a considerable amount of spacing between the grains. This transition structure between the PG and HfC was approximately 10 to 20 µm thick porous layer. The structure was independent of the carbide phase, HfC or TaC, that was deposited. The mixture of hafnium carbide tantalum carbide (HfC/TaC) matrix phase deposited using the same process. Process conditions were selected to produce 4 mol% TaC in the HfC structure. The microstructure of the co-deposited HfC/TaC, including grain size, was similar to HfC (or TaC) microstructures (see Fig. 3). XRD scans of HfC and TaC matched powder diffraction files 39-1491 and 38-1364, respectively. HfC and TaC exhibit very similar X-ray patterns because they have the same cubic cell structure [33, 34] and they have similar lattice parameters (lattice spacings of HfC and TaC  5997  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 1 SEM micrograph of  thick HfC layer  (>120 µm).  (a) Columnar growth showing features that are typical  for high deposition rates. The  compliant transition region between PG and HfC contains voids and enables the growth of a thick HfC layer. (b) Surface morphology of HfC without  any void.  Figure 2 SEM micrograph of as-deposited 30 µm thick TaC coating on Poco ZXF graphite: (a) Surface morphology, (b) Higher magnification, and  (c) Fractured edge of TaC coating showing columnar structure.  5998  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 3  (a) Co-deposited HfC/TaC fracture surface morphology. (b) Higher magnification of the fractured edge showing columnar structure.  Figure 4 Two-dimensional diffraction pattern of HfC and PG. Large Bragg angle (2\\x01) and the azimutal-angle (χ ) measured simultaneously. The  insert, intensity versus 2\\x01, has been obtained by the integration of the data. PG has wide and homogeneous intensity distribution bands (Debye rings).  HfC has discontinuously distributed and rather narrow Debye rings.  5999  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  are 4.64 and 4.0 ˚A, respectively). The two-dimensional images of X-ray diffracting cones intersecting the area detector revealed much less continuity of the diffraction peak for HfC/TaC indicating slightly finer grain structure. The combination of XRD and SEM analyses indicated that a heterogeneous microstructure consisting of HfC matrix and TaC precipitates was produced. The amount of Ta dissolved in the HfC matrix and quan titative phase amounts of HfC and TaC were not determined. The TaC addition (co-deposition of HfC/TaC) did not reduce the grain size and material development focused on the carbon fiber reinforced HfC matrix. Figures 4 and 5 reveal some microstructural information concerning a Cfiber /PGinterface /HfCmatrix composite. An extensive SEM analysis on several fibers revealed that the PC coating was very uniformly coated on the  Figure 5  (a) Fracture morphology of Cfiber /PGinterface /HfC composite. (b) Outer surface morphology of HfC. (c) Columnar structure of HfC, crystal at the start of the growth. (d) Low density HfC layer. (e) PG layer. (f) The edge of fiber fracture surface.  6000  \\x0c', 'fiber fabrics and had very uniform surface texture. The Debye rings obtained from PG are associated with the basal and prism planes of PG and are homogeneously distributed as large bright bands over the azimutalangle. Pronounced individual spots with incomplete Debye rings were correlated with the structures shown in composite Fig. 5b, c and d. The nominal PG coating thickness was approximately 5 µm and some composites with 10 µm PG coating was produced around the fibers without bridging to the neighboring fiber. The thick coating of PG was selected to reduce the heat transfer from the exterior to the fibers during ultra high temperature use. A necessary condition for achieving of high toughness in fiber-reinforced ceramic composites is the promotion of crack deflection and delamination mechanisms. In the present case, this was attained using a lowmodulus interface, PG, between the carbon fiber and the HfC matrix. The micrographs in Fig. 5 illustrates the results for Cfiber /PGinterface /HfCmatrix composite produced by different microstructures at different regions of the constituent phases, (fiber, interface and matrix). The nominal volume fraction of the fiber content estimated to be 20 vol%. The woven fiber cloths were coated with PG to achieve a low modulus, laminated interface between the HfC-matrix and the carbon fiber bundles. The deposited PG, however, exhibited high density and a high degree of preferred orientation of the graphite crystallites as determined by texture analysis of X-ray results (representative example is shown Fig. 4) and caused strain due to thermal expansion mismatch between the PG and the HfC. The porous HfC structure, Fig. 5d, between highly textured PG and fully dense HfC matrix was effective to produce a crack free HfC matrix. The bond strength perpendicular to the deposition plane of PG is very low and readily promotes strain energy release through the frictional dissipation of strain energy, thereby producing a high toughness material. The stress vs. strain behavior of Cfiber /PGinterface / composites was evaluated in tension. Figures 6 and 7 show stress—strain curves for the samples  HfCmatrix  Figure 6 Stress vs. strain behavior of a Cfiber /PGinterface /HfC composite at room temperature. The test conditions are included in the fig ure. This  represents  lowest measured fracture strain capability for a  Cfiber /PGinterface /HfC composite.  ULTRA-HIGH TEMPERATURE CERAMICS  Figure 7 Stress vs. strain behavior of a Cfiber /PGinterface /HfCmatrix composite at room temperature. The test conditions are included in the figure. The graph is representative of a composite with fracture strain (2%). The interrupted data at 1.4% is due to incremental fractures of the composite and accompanying intense deflections of the flags.  which had the lowest and highest strain-to-failure, respectively. High fracture strains (1-2%) were achieved, but the fracture strengths were relatively low. The tensile strength of 30 coupons were evaluated and the average and standard deviation values were 25 and ±8 MPa, respectively.     Cfiber /PGinterface /HfCmatrix  The Cfiber /PGinterface /HfCmatrix  composite was annealed at 2200 C for 4 h in argon. The structural stability was confirmed by the observation that tensile strength for the annealed composites (average 26 and ±6 MPa standard deviation) was the same (within experimental error) as reported above for the as-produced composite. In addition, no obvious structural changes were observed by SEM. As-produced and annealed composites failed between the plies and not at the interface within the porous HfC region. A portion of Cfiber /PGinterface /HfCmatrix composite remained well adhered on the mounting tabs after testing. The fracture surfaces showed a very rough morphology and extensive fiber pull-out was observed. The porous HfC layer (Fig. 5d), located between the PG layer and the fully dense HfC matrix, presumably provided additional mechanism for crack deflection and enhanced strain-to-failure capability. However, the relative contributions of the porous HfC region and the PG interfacial layer for the strain capability are not known. SEM analysis indicated that the PG interface was the primary region for crack deflection. The tensile strength composite is moderate indicating only a small fraction of the fiber strength was available to carry the load. The SEM analysis was not sufficient to identify the weakening mechanism for the composite. It is possible that further engineering is required for the PG interface coating between the carbon fibers and HfC matrix. The present investigation used Rayon fibers, which may be partly responsible for low tensile strength values. Therefore, additional efforts are necessary to extend the present results to the fabrication of composites with alternate carbon fiber types and weave geometries, such  of Cfiber /PGinterface /HfCmatrix  6001  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  as 3D woven structures. 3D woven structures could offer higher strength to improve the shear component of the whole composite and hence is expected to increase the mechanical properties in off-axis loading directions. Additional load bearing capability of the composite could possibly be attained by producing very fine and equiaxed grains of HfC matrix phase.  4. Conclusions  A wide range of HfC, TaC and HfC/TaC coating morphologies on graphite and pyrolytic graphite substrates have been studied and correlated with mechanical properties. The coating adherence to the pyrolytic graphite substrate is achieved using deposits of HfC consisting of randomly oriented fine grains near the substrate. An intermediate region (20 to 30 µm) consisting of porous structure form a ‘compliant’ composite, i.e., a composite with relatively high strain-to-failure. The randomly oriented fine grains grew into large columnar grains and produced dense carbide coatings. Small additions of TaC to the HfC matrix did not result in any obvious change in microstructure compared to HfC samples alone. The necessary requirement for the toughening has been attained using high-strength carbon fibers and the engineering of the fiber matrix interface. Pyrolytic graphite (PG) as a compliant interface ensured crack deflection and its layered structure was the source of strain energy release through frictional dissipation of crack propagation. Further engineering is necessary to reduce PG layer thickness and increase the strain capability. The tensile strength of Cfiber /PGinterface /HfCmatrix composite was 26 MPa (with standard deviation of 8 MPa). The strain-to-failure values for the Cfiber /PGinterface /HfCmatrix composites were larger than 1% and, in many cases, reached as high as 2%. This strain capability exceeds most ceramicmatrix composites, but the tensile strength values are low. Superior fibers with 3D architecture are expected to increase the load bearing capability of the composite. Hafnium carbide was thermodynamically stable in the reducing environment upon heating at 2200 C for four hours. This annealing condition did not change the microstructure, no phase transformations or reactions were observed by SEM and no strength degradation occurred. A viable HfC matrix for carbon reinforced composite for aerospace applications must be “prime reliant” (i.e., it must guarantee protection of a component that would not catastrophically fail or oxidize under use conditions). Additional work is necessary to define the oxidation resistance of composite for time, temperature and stress regimes in oxidizing environment.     the Cfiber /PGinterface /HfCmatrix  References  1. L . Refractory Materials under High Velocity Atmospheric Flight Con K A U F M A N and H .  N E S O R , “Stability Characterization of  ditions,” AFML-TR-69-84, Part II, Vol. II: Facilities and Techniques  Employed for Cold Gas/Hot Wall Tests, ManLabs, Inc., Cambridge,  Mass. (Sept. 1969).  2.  Idem., “Stability Characterization of Refractory Materials under  High Velocity Atmospheric Flight Conditions,” AFML-TR-69-84,  Part  II, Vol.  III: Facilities  and Techniques Employed for Cold  Gas/Hot Wall Tests, MansLabs,  Inc., Cambridge, Mass.  (Sept.  1969).  3.  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ULTRA-HIGH TEMPERATURE CERAMICS  Idem., Surf. Coat. Techn. 36 (1988) 111.  30. 31. C . and  B . T .  B A R G E R O N , P H I L L I P S ,  E .  R .  C .  B E N S O N ,  A .  N .  J E T T E  J. Amer. Ceram.  Soc.  76(4)  (1993)  G .  E M I G ,  S C H O C H and O . W O R M E R , J. de Physique  1040. 32. G . IV(3) (1993) 535. 33. L . E . T O T H , “Transition Metal Carbides and Nitrides” (Academic Press, New York, 1971). 34. E . New York, 1967).  S T O R M S , “The Refractory Carbides” (Academic Press,  K .  Received 9 March 2004 and accepted 20 April 2004  6003  \\x0c', '\\x0c']"
},{
  "_id": 25,
  "PDF": "Changed oxidation behavior of ZrB2–SiC ceramics with the addition of ZrC.pdf",
  "Text": "[\"Available online at www.sciencedirect.com  Ceramics International 41 (2015) 8247-8251  Short communication  CERAMICS  INTERNATIONAL  www.elsevier.com/locate/ceramint  Changed oxidation behavior of ZrB2-SiC ceramics with the addition of ZrC  Hu-Lin Liua,b, Ji-Xuan Liua, Hai-Tao Liua, Guo-Jun Zhanga,n  aState Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai bUniversity of  the Chinese Academy of Sciences, Beijing 100049, China  Institute of Ceramics, Shanghai 200050, China  Received 12 January 2015;  received in revised form 27 February 2015; accepted 27 February 2015  Available online 17 March 2015  Abstract  Oxidation behavior of ZrB2-SiC ceramics containing 0, 10 and 20 vol% ZrC was studied at 1600 1C in stagnant air. The structures of oxide scales depended on ZrC content. The thickness of the SiO2-rich glass layer on the outermost surface decreased steeply with the addition of ZrC. One “ZrC-corroded layer”, containing ZrCxOy, SiC and ZrB2, was observed in ZrB2-SiC-ZrC. The oxidation kinetics was converted from parabolic to linear for the ceramics with 20 vol% ZrC.  & 2015 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  Keywords: Zirconium diboride (ZrB2); Static oxidation; Kinetics; ZrC-corroded layer  1.  Introduction  Transition metal carbides (MeC) such as ZrC, WC and VC have been introduced to ZrB2-SiC ceramics in order to tailor their microstructures and mechanical properties in recent years [1-3]. In these ternary ceramic systems, ZrC phase always existed due to the reaction and solid solution between MeC and other phases [4]. Therefore, the effects of ZrC on properties of ZrB2-based ceramics attract researchers' attentions, for example, the oxidation behavior of ZrB-SiC containing ZrC (ZSZ). Guo et al. [5] reported oxide structure of ZrB2-SiC with addition of 15 vol% ZrC at 1500 1C and observed one ZrO2 layer with large voids. This layer could not effectively prevent oxygen diffusion, compared with the borosilicate glass layer on the surface of ZrB2-SiC [6-8]. Wu and Wang et al. [9,10] studied the oxidation behavior of ZrB2-SiC containing 6 vol% ZrC at 1600 1C and 1750 1C, respectively. The layered structures consisted of one SiO2-rich glass layer, one thin layer of ZrO2-SiO2 and one SiC-depleted layer containing ZrO2 and ZrB2. Zhang et al. [11] investigated oxidation kinetics of ZrB2-SiC incorporating 40 vol% ZrC at 1200-1500 1C. They pointed out that the oxidation was reaction controlled below  nCorresponding author. Tel.: þ 86 21 52411080; fax: þ 86 21 52413122. E-mail address: gjzhang@mail.sic.ac.cn (G.-J. Zhang).  http://dx.doi.org/10.1016/j.ceramint.2015.02.150  0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  1200 1C, and changed into diffusion controlled above 1300 1C. However, there is limited work focused on the influence of content of ZrC on oxidation kinetics and structures of oxide scales in ZSZ ceramics despite the importance of chemical composition. To fill the gap and provide in-depth understanding of the characteristics of ZSZ ceramics, the effects of composition on oxidation behavior of ZSZ will be investigated in this paper. Static oxidation experiments for specimens containing various content of SiC (10 and 20 vol%) and ZrC (0, 10 and 20 vol%) will be performed in air at 1600 1C with dwelling time up to 4 h. Microstructural features after oxidation will be examined and discussed in details.  2. Experimental procedure  The raw materials were home synthesized ZrB2 powder [12] (D50 ¼ 1.05 μm, 98% purity), α-SiC powder (D50 ¼ 0.45 μm, 98.5% purity, Changle Xinyuan Carborundum Co. Ltd., (D50 ¼ 0.85 μm, 99% purity). The chemical composition (vol Shandong, China) and home synthesized ZrC powder [13] %) for four batches of powder mixtures are as follows: ZrB2- 20SiC, ZrB2-10SiC-10ZrC, ZrB2-20ZrC and ZrB2-20SiC- 20ZrC (marked as ZS20, ZS10Z10, ZZ20 and ZS20Z20).  \\x0c\", '8248  H.-L. Liu et al.  / Ceramics International 41 (2015) 8247-8251  Fig. 1. The formation of the “ZrC-corroded layer” in ZS20Z20 after oxidized for 2 h: (a) the whole oxide scale, (b) and (c) higher magnification images of the white  and black rectangle in (a), and (d) EDS spectra in different  regions in (b).  Powder mixtures were ball-milled in alcohol for 24 h in polyethylene jars using Si3N4 milling medium balls, and dried by rotary evaporation. Then the mixed powder compacts were hot pressed at 1900 1C and 30 MPa for 1 h. The furnace was heated under vacuum below 1600 1C. Above this temperature, the atmosphere was switched to flowing argon gas. billets with dimensions of 3 \\x02 4 \\x02 15 mm3, and all The bars for oxidation testing were cut from the sintered the surfaces were polished to 0.5 μm using diamond paste. Before oxidation, they were cleaned in an ultrasonic bath with alcohol. After dried, specimens were placed on a zirconia support with ridges to minimize the contact area. And oxidation tests were conducted in a muffle furnace. Specimens were heated at 5 1C/ min to 1600 1C and held for 0, 0.5, 1, 2 or 4 h in stagnant air. The mass of specimens before and after oxidation was measured using a balance with 0.1 mg precision. Normalized mass gain (mg/cm2) was calculated from mass gain and surface area which was calculated based on the measured dimensions of specimens. The microstructures of the oxidized specimens were characterized using scanning electron microscopy (SEM, TM3000, Hitachi, Tokyo, Japan) with energy-dispersive spectroscopy (EDS). And the thicknesses of oxide scales were measured using SEM images of polished cross sections.  3. Results and discussion  ZrC exhibits different oxidation behavior compared with SiC. The main reactions of SiC and ZrC with O2 can be expressed as follows [14,15]: SiC (s) þ 3/2O2 (g)-SiO2 (l) þ CO (g)  (1)  SiC (s) þ O2 (g)-SiO (g) þ CO (g) ZrC (s) þ 3/2O2 (g)-ZrO2 (s) þ CO (g) ZrC (s) þ O2 (g)-ZrCxOy (s) þ CO (g)  (2)  (3)  (4)  The Gibbs energy ΔG0 at 1600 1C for \\x00 797.168, \\x00 252.142 and \\x00 842.910 kJ mol-1, respectively, using reactions (1)-(3) was data from the NIST-JANAF tables [16]. Although no thermodynamic data were available ZrCxOy, ΔG0 of reaction (4) would be smaller than reaction (3) considering reaction (4) was partial oxidation of ZrC. It indicates the oxidation of ZrC is easier than SiC based on thermodynamics. Combining ΔG0 with equilibrium constant Keq by Eq. (5), the reaction having lower Gibbs energy could occur under lower oxygen pressure. In other words, ZrC could be partial oxidized by reaction (4) under oxygen partial pressure that SiC was stable enough to avoid oxidation.  ΔG0 ¼ \\x00 RT ln K eq  ð5Þ  where R is the ideal gas constant and T is the absolute temperature. It has been reported that SiC could be oxidized by reaction (2) under oxygen pressure that ZrB2 was stable [17]. Therefore, the equilibrium oxygen pressures for the oxidation of these three order: ZrB24 SiC4 ZrC. phases should be in the following Considering oxygen partial pressure decreased from the surface to the unreacted bulk during oxidation, a region called a SiCdepleted layer existed [17], in which SiC and ZrC were oxidized by reactions (2) and (3). And a region, that only ZrC was oxidized, should also exist in ZrB2-SiC-ZrC. Fig. 1(a) shows the SEM image of the polished section of ZS20Z20 after oxidized for 2 h. The microstructures of the region behind SiC-depleted layer (the white rectangle) and unreacted bulk (the black rectangle) were different as shown in  \\x0c', 'H.-L. Liu et al.  / Ceramics International 41 (2015) 8247-8251  8249  Fig. 2. Typical SEM images of  the polished sections of different ceramics after oxidized for 2 h:  (a) ZS20,  (b) ZS20Z10, and (c) ZS20Z20.  Fig. 1(b) and (c). In Fig. 1(c), the gray, black and white phases are ZrB2, SiC and ZrC, respectively. However, some small particle clusters containing pores are observed in Fig. 1(b), instead of white phase. These clusters contained Zr, C and O according to the EDS analysis. They should be ZrCxOy, not the mixtures of ZrO2 and ZrC, because of the uniform oxidation of ZrC grains. Furthermore, ZrB2 and SiC also existed in this oxidized region as shown in Fig. 1(c). These experimental results were consistent with the thermodynamic analysis. Here we name this oxidized layer as a “ZrC-corroded layer”. Oxide scale structures were also influence by the ZrC content. Fig. 2 shows typical SEM images of the polished sections of different ceramics after oxidized for 2 h. The oxide scale of ZS20 consisted of three layers: (1) one SiO2-rich layer (52 μm); (19 μm); and (3) one SiC(2) one ZrO2-SiO2 layer (48 μm), which is depleted layer in agreement with the previous studies [18]. For the ceramic containing 10 vol% ZrC, the SiO2-rich glass layer became much thinner (15 μm) and an extended ZrC-corroded layer (306 μm) was formed behind the SiC-depleted layer (65 μm). The thickness of the ZrO2-SiO2 layer was 65 μm. However, the outer SiO2-rich layer disappeared when the ZrC content increased to 20 vol%. And the thicknesses of ZrO2-SiO2 layer and SiC-depleted layer were 224 μm and 85 μm. In addition, large pores and cracks were also observed in ZrO2-SiO2 layer, as a result of the upwelling of aggregate CO gas. These results indicated that the added ZrC was harmful to the oxidation resistance of ZrB2-based ceramics. Fig. 3(a) presents the thickness of the SiO2-rich layer versus the oxidation time in different ceramics. The thickness reduced with the increasing ZrC content, and this oxide layer disappeared when the ZrC content was above 20 vol%. The thickness was between 40 μm and 70 μm for ZS20 when the dwelling time was between 0.5 h and 4 h, while it was between 5 μm and 20 μm for the ceramic with 10 vol% ZrC addition. And the thickness changed slightly with SiC content for the ceramics containing the same amount of ZrC. It indicated that the SiO2rich layer was controlled by the ZrC content. The thickness of SiO2-rich layer determined the oxidation kinetics. The kinetics was analyzed based on mass gain varying with the dwelling time, as shown in Fig. 3(b). A typical power rate equation was utilized in this work [19,20]:  mn ¼ k t  ð6Þ  where m is  the  change  in mass gain, n is  the  exponent,  k  is  the rate constant and t  is the oxidation dwelling time. In addition,  mass gain of ZSZ could not be neglected during the temperature rise period according to Fig. 3(b). Therefore,  the value of m at a  given dwelling time should subtract mass gain at  the dwelling  time of 0 h. After data fitting, we obtained the values of n for  ZS20, ZS20Z10, ZS10Z10 and ZS20Z20, which were 1.96, 1.38,  1.13 and 0.98,  respectively. The oxidation was converted from  parabolic kinetics for ZS20 to linear kinetics for ZS20Z20. These  kinetic  characteristics were  consistent with  the  oxide  scale  structures. When the ZrC content  in the  ceramics was below  10 vol%,  the protective SiO2 layer on the ceramic surface led to the parabolic kinetics, indicating a diffusion controlled mechan ism. For  the  ceramic  layer vanished and the  containing 20 vol% ZrC,  the outer SiO2 linear oxidation occurred, indicating a  reaction controlled mechanism. The large volume change of ZrC during the oxidation and the formation process of SiO2 liquid should determine the oxide structure and thickness of each oxide layer. Oxidation of both ZrB2 and SiC were negligible below about 800 1C [14]. the obvious oxidation of ZrC began above 400 1C In contrast, [21,22]. During 400-800 1C, only ZrC phase was oxidized to ZrO2 with a large volume expansion (32%). In this stage, the formed ZrO2 grew vertically to the specimen surface due to the space limitation by the unreacted matrix. And the ZrO2 layer extended beyond the outer surface of the specimen, as shown schematically in Fig. 4(b). In addition, CO gas was also produced from the oxidation of ZrC, when ZrB2 phase was partly substituted by ZrC. Channels must exist to release CO gas. These channels could also serve as the oxygen diffusion path, which resulted in the further oxidation of ZrC to ZrO2. 1200 1C, ZrB2 At higher temperature between 800 and began to be oxidized with the formation of ZrO2 (with 9% volume expansion) and B2O3 liquid. In this stage, both ZrB2 and ZrC were oxidized and ZrO2 skeleton was filled with B2O3, which led to the different oxide structure indicated by Fig. 4(c). O2 would diffuse through the formed ZrO2-B2O3 layer and the ZrO2 volume further increased. SiC particles did not be oxidized significantly and incorporated into the oxide scale. With the decreasing oxygen pressure beneath the outer ZrO2-B2O3 layer, ZrC would convert to ZrCxOy. 1200 1C, When ZrB2-SiC-ZrC was heated above the composition and structure of the oxide scale was changed,  \\x0c', '8250  H.-L. Liu et al.  / Ceramics International 41 (2015) 8247-8251  Fig. 3.  (a) The thickness of  the SiO2-rich layer and (b) oxidation mass gain versus the oxidation dwelling time.  Fig. 4. Schematic diagram of the oxidation process for ZrB2-SiC-ZrC in stagnant air with elevated temperatures: (a) unoxidized bulk, (b) initial response that only ZrC was oxidized during heating (400-800 1C), (c) oxidation of ZrB2 and ZrC in the middle temperature (800-1200 1C), and (d) steady state at high temperature (1200-1600 1C).  due to the evaporation of B2O3 and oxidation of SiC. The ZrO2-SiO2 layer replaced the ZrO2-B2O3 layer formed below 1200 1C, as shown in Fig. 4(d). The SiO2-rich layer was formed at the surface. However, because of the large porosity of the formed ZrO2, SiC was oxidized in situ by reaction (1). The formed SiO2 glass filled in ZrO2 skeleton. Simultaneously, the SiO gas formed from SiC-depleted layer under a lower oxygen partial pressure diffused into ZrO2-SiO2 layer and reacted with O2 by reaction (7). As the larger porosity in ZrO2 resulted from added ZrC, the SiO was able to be oxidized before transporting into the outer SiO2-rich layer. The newly formed SiO2 was entrapped in ZrO2-SiO2 layer. Therefore, the ZrO2-SiO2 layer in ZSZ is thicker than in ZS and the thickness of SiO2-rich layer was much thinner than ZrO2-SiO2 layer in ZSZ. Based on the experimental results, the SiO2-rich layer in ZSZ was vanished when the ZrC volume was above 20%, compared with ZS.  SiO (g) þ 1/2O2 (g)-SiO2 (l)  (7)  In addition, although the added ZrC would reduce static oxidation resistance due to the thinner SiO2-rich layer, it played different roles on the ablation resistance in combined conditions of high temperatures, high-speed gas flow and low oxygen partial pressure [23]. The outmost SiO2-rich layer in ZS will be eroded by shear forces derived from gas flow during ablation process [24]. However, ZrO2-SiO2 layer became thicker when ZrC was introduced into ZrB2-SiC, ZrO2 solid phase would increase the viscosity and strength of this layer compared with SiO2 glass layer. Therefore, this layer was stronger to resist mechanical erosion. We have preliminarily studied the ablation resistance of ZrB2-20SiC and ZrB2- 20SiC-5ZrC (number in vol%) at 1900 1C using oxyacetylene torch testing. Spallation was observed on the surface of ZrB2- 20SiC after ablation, whereas the oxide layer covered on the  \\x0c', 'H.-L. Liu et al.  / Ceramics International 41 (2015) 8247-8251  8251  surface of ZrB2-20SiC-5ZrC without obvious spallation, pores or cracks. It indicated that the added ZrC improved the ablation resistance of ZrB2-SiC. The details of ablation behavior will be investigated and reported in the further.  4. Conclusions  In summary, the static oxidation behavior of ZrB2-SiC ceramics as a function of ZrC content was investigated at 1600 1C. The addition of ZrC played a dominant role on the oxidation resistance of ZrB2-SiC-ZrC. The SiO2-rich layer was much thinner with 10 vol% ZrC (5-20 μm) and disappeared at 20 vol% ZrC. One ZrC-corroded layer, containing ZrCxOy, ZrB2 and SiC, was observed in ZSZ after oxidation. The obvious static oxidation of ZSZ was attributed to the non-protective oxidation of ZrC above 400 1C and the large volume expansion during the conversion of ZrC to ZrO2 (about 32%). The oxidation of baseline ZS and low content of ZrC (10 vol%) obeyed parabolic kinetics while that containing 20 vol% ZrC followed linear kinetics.  Acknowledgments  The authors express their gratitude to financial supports from the National Natural Science Foundation of China (No. 51272266), the State Key Laboratory of High Performance Ceramics and Superfine Microstructure of Shanghai Institute of Ceramics are gratefully acknowledged.  References  [1] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Pressureless densification of  zirconium diboride with boron carbide additions, J. Am. Ceram. Soc. 89  (2006) 1544-1550.  [2]  J.  Zou, G.J.  Zhang, Y.M. Kan,  ceramics with VC addition:  P.L. Wang, Hot-pressed ZrB2-SiC chemical reactions, microstructures, and  mechanical properties, J. Am. Ceram. 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Zhang, Oxidation resistance and strength retention of  ZrB2-SiC ceramics, J. Eur. Ceram. Soc. 30 (2010) 2387-2395. [20] C.M. Chen, L.T. Zhang, W.C. Zhou, M.Q. Li, High temperature  oxidation of LaB6-ZrB2 eutectic in situ composite, Acta Mater. 47 (1999) 1945-1952.  [21] S. Shimada, T. Ishil, Oxidation kinetics of zirconium carbide at relatively  low temperatures, J. Am. Ceram. Soc. 73 (1990) 2804-2808.  [22] S. Shimada, A thermoanalytical study on the oxidation of ZrC and HfC  powders with formation of carbon, Solid State Ion. 149 (2002) 319-326.  [23] H. Wu, C. Xie, W. Zhang, J. Zhang, Fabrication and properties of 2D C/  C-ZrB2-ZrC-SiC composites by hybrid precursor olysis, Adv. Appl. Ceram. 112 (2013) 366-373.  infiltration and pyr [24] X. Zhang, P. Hu, J. Han, S. Meng, Ablation behavior of ZrB2-SiC ultra high temperature ceramics under simulated atmospheric re-entry condi tions, Compos, Sci. Technol. 68 (2008) 1718-1726.  \\x0c']"
},{
  "_id": 26,
  "PDF": "Characterization of hot-pressed short carbon fiber reinforced ZrB2–SiC ultra-high temperature ceramic composites.pdf",
  "Text": "['Journal of Alloys and Compounds 472 (2009) 395-399  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j a l l c o m  Characterization of hot-pressed short carbon fiber reinforced ZrB2-SiC ultra-high temperature ceramic composites  Feiyu Yang ∗ , Xinghong Zhang, Jiecai Han, Shanyi Du  Center for Composite Materials and Structure, Harbin Institute of Technology, Harbin 150001, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 7 February 2008 Received in revised form 18 April 2008 Accepted 22 April 2008 Available online 10 June 2008  Keywords:  ZrB2 -SiC ceramics Short carbon fiber Microstructure Mechanical properties Oxidation  1.  Introduction  Densified ZrB2 -SiC-based ultra-high temperature ceramics reinforced with short carbon fibers were prepared by conventional hot-pressing. The microstructure, mechanical and oxidation resistance properties of the composite were investigated. The Csf /ZrB2 -SiC composite had improved fracture toughness of 6.6 MPa m1/2 compared to ZrB2 -SiC composite of 4.3 MPa m1/2 due to fiber debonding, fiber pull-out and fiber bridging as well as crack deflection. It was found that the low modulus of carbon fiber and a graphitization transition layer between fiber and matrix led to the decreased flexural strength. The oxidation resistance tests were carried out on Csf /ZrB2 -SiC using an oxyacetylene torch. The temperature of the oxidized specimen exceeded 1800   C and the surface layer appeared dense and adherent. No macro-cracks or spallation were detected, suggesting that these composites possess a better oxidation resistance than ZrB2 -SiC. The improved oxidation resistance is attributed to the formation of a coherent SiO2 rich scale, which acts as an effective barrier against the inward diffusion of oxygen. The results presented here point to a potential way for improving toughness of composite without sacrificing oxidation protection. © 2008 Elsevier B.V. All rights reserved.  Most modern designs of hypersonic vehicles incorporate sharp aerosurfaces to increase aerodynamic performance. These designs require materials capable of operating in extreme reentering environment, such as oxidizing atmosphere at high temperatures (above 1500  C) and corrosive gases at high velocities. Zirconium diboride (ZrB2 ), which is referred to as ultra-high temperature ceramics (UHTCs) have been proposed as candidates for these applications due to its high melting temperature (3040  C), good oxidation resistance and thermo-mechanical properties. When combined with SiC, ZrB2 -based composites exhibit indeed excellent refractoriness, high oxidation resistance and are as such good potential candidates for the above-mentioned applications. So ZrB2 -SiC is currently considered the baseline UHTCs [1-5]. Nonetheless its low fracture toughness has long prevented this material from being used for wide application. Its susceptibility to brittle fracture can lead to unexpected catastrophic failure. One major research direction has been to increase its fracture toughness by incorporating fiber into the base material to form a ceramic matrix composite due to their reduced weight and damage tolerant behavior [6-8].  Several processes, including tape casting/lamination, chemical vapour infiltration (CVI), solid and liquid infiltration followed by hot pressing (HP) or hot isostatic pressing (HIP) have been employed for the fabrication of continuous fiber reinforced materials. Meanwhile time consuming decrease and significant decrease in the manufacturing cost of components are the driving forces for simple and rapid way to be developed by a cost-effective method. This can be partly alleviated by use of short carbon fibers. These short fibers are characterized not only by lower price and good eligible mechanical properties but also by the economic processing methods compared to long fiber reinforcement. Using short fibers, the composites can be manufactured by conventional production methods like hot pressing, injection molding or extrusion. Besides short carbon fibers lead to isotropic properties of the composite because of the homogenous and three-dimensional reinforcement [9-12]. Meanwhile the incorporation of short fibers introduces new problems such as the oxidation of reinforcing phase and the interface between fibers and matrix. The purpose of this work is to report the processing method and evolution way of the microstructure, mechanical and oxidation resistance properties of short carbon fiber reinforced ZrB2 -SiC composite (abbreviated as Csf /ZrB2 -SiC).  2.  Experimental procedure  ∗  Corresponding author. Tel.: +86 451 86402382; fax: +86 451 86402382. E-mail address: yangfyhit@sina.com.cn (F. Yang).  Commercially available raw powders were used in this study. The ZrB2 and SiC powders (Northwest Institute for Non-Ferrous Metal Research, China) had reported  0925-8388/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.04.092  \\x0c', '396  F. Yang et al. / Journal of Alloys and Compounds 472 (2009) 395-399  Fig. 1. SEM micrograph from a polished cross-section of (a) ZrB2 -SiC and (b) Csf /ZrB2 -SiC composite.  purity of 99% and 98.7%, and particles size of 2.5 \\u242em and 1 \\u242em, respectively. The carbon fiber selected for this study was T800H fiber (Torayca Co., Ltd., Tokyo, Japan) with tensile strength and elastic modulus of about 5490 MPa and 294 GPa, respectively. The T800H fiber had a diameter of 5 \\u242em and it was chopped into 2 mm short fibers. The powder mixtures ZrB2 plus 20 vol.% of SiC and 20 vol.% of short carbon fibers were ball-milled for 10 h in a polyethylene bottle using ZrO2 balls and ethanol as the grinding media. To minimize segregation by sedimentation during drying, solvent removal was performed using rotary evaporation. Then the dried powder mixtures were directly hot-pressed in a boron nitride coated graphite die at 2000   C for 60 min under vacuum (50 Pa), 30 MPa. Specimens were polished and ultrasonic cleaned. The resulting microstructures of polished or fractured sections of composites were studied by scanning electron microscopy (SEM, FEI Sirion, Holland) with simultaneous chemical analysis by energy dispersive spectroscopy (EDS, EDAX Inc., USA). The bulk density of the hot-pressed billets was determined using the Archimedes technique with water as the immersing medium. Flexural strength in a three-point configuration was tested on 3 mm × 4 mm × 36 mm chamfered bar, using a 30 mm span and a crosshead speed of 0.5 mm min−1 . Fracture toughness (KIC ) was evaluated by a single-edge notched beam test with a 16 mm span and a crosshead speed of 0.05 mm min−1 using 2 mm × 4 mm × 22 mm test bars, on the same jig used for the flexural strength. A Vickers indenter with 9.8 N as applied load for 15 s on polished sections was used to measure the micro-hardness (Hv 1.0) and observe the crack propagation. To evaluate the interface structure of the composites, specimen for transmission One Csf /ZrB2 -SiC sample coupon of ˚ 19 mm × 14 mm was cut from hot-pressed electron microscopy (TEM) was examined in detail by JEM-2010 (TEM, JEOL, Japan). plates for oxidation tests. Oxidation tests were carried out with an oxyacetylene torch according to GJB323A-96 standard on the oxyacetylene ablation equipment [13]. Models were placed in graphite holders which enabled test durations in excess of 180s. The pressure and flux of acetylene were 0.1 MPa and 1.15 m3 h−1 , and for oxygen 0.5 MPa and 2.30 m3 h−1 , respectively. The emissivity was determined by a two-color radiation pyrometer and the temperature of specimen surface was measured using a one-color pyrometer (Mikron Instrument Co., Inc., Oakland, NJ). The Tested specimens were photographed, and X-ray diffraction (XRD; Rigaku, Japan) and SEM characterizations were performed.  3. Results and discussion  3.1. Microstructure  Fig. 1 showed scanning electron micrograph of the polished surface of the ZrB2 -SiC and Csf /ZrB2 -SiC composite. Through the analysis of EDS (not shown here), the SiC particle was found well dispersed in the ZrB2 matrix (Fig. 1a), and a uniform distribution  Table 1 Mechanical properties of Csf /ZrB2 -SiC and ZrB2 -SiC composite (measured density (cid:2), fracture toughness KIC , flexural strength (cid:3) , micro-hardness Hv 1.0)  Sample  (cid:2) (g/cm3 )  Relative density (%)  Csf /ZrB2 -SiC ZrB2 -SiC  4.63 5.50  99.3 99.8  KIC (MPa m1/2 )  (cid:3) (MPa)  Hv 1.0 (GPa)  6.6 ± 0.1 4.3 ± 0.2  445 ± 36 502 ± 45  19.2 ± 0.8 20.7 ± 0.9  of short fibers and the SiC particulate in ZrB2 matrix (Fig. 1b) was detected. For Csf /ZrB2 -SiC material, the microstructure was regular and no agglomeration was observed. Although raw carbon fibers had a high aspect ratio and they were prone to be twisted during mixing, the wet ball-milling technique could effectively prevent fibers from aggregating. When a combination of density measurements (Table 1) and SEM analysis was used, no indication of porosity was found in ZrB2 -SiC or Csf /ZrB2 -SiC billets (relative density >99%), which showed that hot-pressing way was effective for the densification of ZrB2 -SiC or Csf /ZrB2 -SiC composite.  3.2. Mechanical properties  Results of the mechanical properties were listed in Table 1. The measured fracture toughness of Csf /ZrB2 -SiC was 6.6 MPa m1/2 , which increased by approximately 54% for ZrB2 -SiC of 4.3 MPa m1/2 . The improvement of fracture toughness was due to the short fiber reinforcement. Fig. 2 showed the fractured surface of specimens for testing fracture toughness. On the fractured surface of Csf /ZrB2 -SiC composite (Fig. 2b), some short fibers fractured were detected, some short fibers debonding and pull-out were observed, which showed that during tests the crack overcame not only propagating resistance of matrix, but also overcame the interfacial shear resistance to make the fibers debonding and pull-out. The rough fracture surface of the composite and ragged crack propagation path showed that crack deflection should be another toughening mechanism because the crack swerving and twisting along fiber/matrix interface consumed more energy than  Fig. 2. SEM micrograph of the fracture surface of (a) ZrB2 -SiC and (b) Csf /ZrB2 -SiC composite.  \\x0c', 'F. Yang et al. / Journal of Alloys and Compounds 472 (2009) 395-399  397  Fig. 3. SEM micrograph of indentation-induced crack propagation on a polished surface showing fiber bridging.  Fig. 6. Measured results of temperature curves of the Csf /ZrB2 -SiC sample.  The addition of short fibers also had influence on the strength of composites. Fig. 2 showed that the growth of matrix grains was inhibited due to the existence of short fibers. The ZrB2 grains changed from rod-shape (7 \\u242em in length and 5 \\u242em in width) in to equiaxed structure (3 \\u242em) composite (Fig. 2b), and the grain size of SiC (1 \\u242em) ZrB2 -SiC (Fig. 2a) in Csf /ZrB2 -SiC in both composites did not remarkably increase. The measured flexural strength of Csf /ZrB2 -SiC composite was 445 MPa, which decreased by 11% for ZrB2 -SiC composite (502 MPa). Though the decrease of grain size was beneficial for improving flexural strength, the modulus of fiber (294 GPa) was much lower than the reported modulus of matrix (466 GPa) [14], which was thought to decrease the strength of the composite. Furthermore, it was reported that heat treatment could improve the graphitization degree of carbon fiber [15,16], and for Csf /ZrB2 -SiC composite, it experienced a sintering process up to 2000  C. Fig. 4 showed a TEM image of cross section of one short carbon fiber in Csf /ZrB2 -SiC composite. Evidently, the outside layer of fiber with 2 um in thickness was graphitization transition layer and some platelet structure crystallites with different orientations in this layer were observed. The graphite flake structure in graphitization transition layer was easy to slip, which would result in the relatively weaker bond between fiber and matrix. This inference was in accordance with SEM observations on fracture surface (Fig. 2b), where the fiber debonding and fiber pull-out indicated that the interface bond between fiber and matrix was not strong. Thus the graphitization transition layer of the short fiber surface also led to the decreased flexural strength. Here it is worth noting that for UHTC, the graphitization transition layer can not only help fiber to realize the pull-out toughening mechanism, but also release strain energy by the slip of graphite  Fig. 4. TEM detail of the cross-section of a short carbon fiber in Csf /ZrB2 -SiC composite with surface graphitization (arrowed).  crack propagating directly. In addition, the crack propagation paths at the corners of indents during microhardness tests, was shown in Fig. 3, and the fiber bridging interaction zone behind the crack tip could be seen. It was believed that these interaction effects absorbed crack propagating energy during fracture and led to the improved toughness.  Fig. 5. Photographs of Csf /ZrB2 -SiC (a) before and (b) after oxyacetylene torch testing.  \\x0c', '398  F. Yang et al. / Journal of Alloys and Compounds 472 (2009) 395-399  Fig. 7. Surface morphologies of the oxide scale after Csf /ZrB2 -SiC specimen exposure to oxyacetylene torch testing.  flake structure when subjected to thermal load, resulting in the improved reliability at ultra-high temperature conditions.  3.3. Oxidation resistance of Csf /ZrB2 -SiC composite  The performance of Csf /ZrB2 -SiC was evaluated under oxyacetylene torch condition for 180 s. Photographs of the Csf /ZrB2 -SiC composite before and after oxyacetylene torch testing are shown in Fig. 5. The oxidized surface layer appears compact and adherent. No macro-cracks or spallation were detected there. The surface temperature increased drastically in about 30 s to 1500  C and then began to rise gradually, reaching a maximum temperature of 1890  C when the sample was exposed to the torch heater, as can be seen in Fig. 6. Under such a high thermal stress, no cracks were observed on the sample, however, extensively damaged surface and spalling outer oxide scale of ZrB2 -SiC were observed by Zhang et al. [17] after torch testing under the same conditions. These observa tions suggest that the Csf/ZrB2 -SiC composite may offer improved thermal shock resistance compared to ZrB2 -SiC, but further testing is required to confirm this assertion. Fig. 7 showed the surface morphology of the oxide scale with different magnifications. Some micro-bubbles and craterlets with a few small pores upon the surface were found, which presumably derived from the evolution of volatile products like SiO(g), CO(g), CO2 (g) and B2O3 (g) during oxidation. These features are commonly observed and reported in the UHTC literatures [18,19]. The surface-oxidized sample consisted of two distinctly different scales, namely, the bright phase and gray phase. The dominant bright phase and a small amount gray phase was identified, respectively, as monoclinic ZrO2 and silicacontaining glass (mostly SiO2 ) according to the combination of EDS and XRD examination (not shown). Carbon was not detected by the apparatus, showing that all carbon fibers on the surface have been oxidized. The cross-section analysis (Fig. 8) also revealed the presence of 60 \\u242em oxidized scale for Csf /ZrB2 -SiC, in contrast to the  Fig. 8. SEM result for a cross-section of Csf /ZrB2 -SiC after oxyacetylene torch testing.  \\x0c', 'F. Yang et al. / Journal of Alloys and Compounds 472 (2009) 395-399  399  oxidized scale thickness of 600 \\u242em reported for ZrB2 -SiC [20], which showed the better oxidation resistance of Csf /ZrB2 -SiC composite than ZrB2 -SiC. It was significant to note that, SiO2 rich glass continuously covered and sealed the space between ZrO2 particles. Such a surface structure was responsible for the improvement in oxidation resistance of the composites due to the formation of viscous SiO2 glass which has a higher melting temperature and a lower vapor pressure. For Csf /ZrB2 -SiC, though carbon fibers on the surface were very easy to be oxidized, SiO2 rich glass formed on the surface acted as an effective barrier to oxygen diffusion. The better oxidation resistance also made Csf /ZrB2 -SiC material very potential for service in oxidizing atmosphere at high temperatures.  4. Conclusion  An investigation of the effect of addition of 20 vol.% short carbon fiber on microstructure, mechanical properties and oxidation resistance of hot-pressed ZrB2 -SiC ceramic matrix composite has been conducted. It was observed that short fiber addition can help to enhance the fracture toughness from 4.3 MPa m1/2 for ZrB2 -SiC composite to 6.6 MPa m1/2 for Csf /ZrB2 -SiC composite. The observed toughening mechanisms were attributed to fiber debonding, fiber pull-out, fiber bridging and crack deflection. The flexural strength of Csf /ZrB2 -SiC composite was 445 MPa, a decrease of 11% over that of ZrB2 -SiC composite, which was due to the lower modulus of carbon fiber and the existence of graphitization transition layer between fiber and matrix. The oxidation resistance of Csf /ZrB2 -SiC composite was tested by oxyacetylene torch with a dwell time of 180s and the surface temperature of 1890  C. Short carbon fibers on the surface have been oxidized and a ZrO2 rich oxidized layer containing viscous SiO2 glass was observed. This oxidized layer was believed to be very important for oxidation resistance, as it formed a passive layer that slowed the oxidation.  The results presented in this paper point to a potential method for improving not only the fracture toughness but also the oxidation resistance of ultra-high temperature ceramic.  Acknowledgements  This work was supported by the National Natural Science Foundation of China (90505015 and 50602010), the Research Fund for the Doctoral Program of Higher Education (20060213031) and the Program for New Century Excellent Talents in University.  References  58  (1997)  I.G. Talmy,  J.A. Zaykoski,  J. Am. Ceram. Soc. 90  Soc. Bull.  J. Am. Ceram. Soc. 90 (2007)  J.W. Halloran, C.E. Henderson,  [1] S.N. Karlsdottir, 2863-2867. [2] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, J. Eur. Ceram. Soc. 22 (2002) 2757-2767. [3] C. Mroz, Am. Ceram. Soc. Bull. 73 (1994) 141-142. [4] K. Upadhya, J.M. Yang, W.P. Hoffmann, Am. Ceram. 51-56. [5] W.G. Fahrenholtz, G.E. Hilmas, (2007) 1347-1364. [6] D.B. Marshall, A.G. Evans, J. Am. Ceram. Soc. 68 (1985) 225-231. [7] J.J. Brennan, K.M. Prewo, J. Mater. Sci. 36 (1982) 2371-2393. [8] R.E. Tressler, Composites Part A 30 (1999) 429-437. [9] F. Folgar, Ceram. Eng. Sci. Proc. 9 (1988) 561-578. [10] J.S. Lee, M. Imai, T. Yano, Mater. Sci. Eng. A 339 (2003) 90-95. [11] G.M. Song, Q. Li, G.W. Wen, Y. Zhou, Mater. Sci. Eng. A 326 (2002) 240-248. [12] X.Y. Wang, F. Luo, X.M. Yu, D.M. Zhu, W.C. Zhou, Scripta Mater. 57 (2007) 309-312. [13] G.M. Song, Y. Zhou, Y.J. Wang, Mater. Charact. 50 (2003) 293-303. [14] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Ceram. Soc. 87 (2004) 1170-1172. [15] S. Wang, Z.H. Chen, W.J. Ma, Q.S. Ma, Ceram. Int. 32 (2006) 291-295. [16] G.B. Zheng, H. Sano, Uchiyama, Carbon 41 (2003) 853-856. [17] X.H. Zhang, P. Hu, J.C. Han, L. Xu, S.H. Meng, Scripta Mater. 57 (2007) 1036-1039. [18] W.G. Fahrenholtz, J. Am. Ceram. Soc. 88 (2005) 3509-3512. [19] W.G. Fahrenholtz, J. Am. Ceram. Soc. 90 (2007) 142-148. [20] F.Y. Yang, X.H. Zhang, J.C. Han, S.Y. Du, J. Inorg. Mater. 4 (2008), in press.  \\x0c']"
},{
  "_id": 27,
  "PDF": "Characterization of novel ceramic composites for rocket nozzles in high-temperature harsh environments.pdf",
  "Text": "['International Journal of Heat and Mass Transfer 163 (2020) 120492   Contents lists available at ScienceDirect   International Journal of Heat and Mass Transfer   journal homepage: www.elsevier.com/locate/hmt   Characterization of novel ceramic composites for rocket nozzles in   high-temperature harsh environments   Stefano Mungiguerra  , Giuseppe D. Di Martino  a , Raffaele Savino  a , Luca Zoli  b , Diletta Sciti  Laura Silvestroni   a ,   b ,   ∗  b   a University of Naples “Federico II”, Department of Industrial Engineering, P.le Vincenzo Tecchio 80, 80125 Napoli, Italy   b National Research Council, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy   a r t i c l e   i n f o   a b s t r a c t   Article history:   Received 12 June 2020   Revised 2 September 2020   Accepted 18 September 2020   Keywords:   Ultra-high-temperature ceramic matrix   composites   Hybrid rocket nozzles   Innovative test set-up   Computational fluid dynamic simulation   Thermo-chemical erosion   This paper presents the results of experimental tests for the characterization of Ultra-High-Temperature   Ceramic Matrix Composite (UHTCMC) materials for near-zero erosion rocket nozzles. Two dedicated test   set-ups were developed for preliminary screening of material candidates in a representative environment,   characterized by relevant heat flux and temperature. The experimental set-up was based on a lab-scale   200N-class hybrid rocket engine, employing gaseous oxygen as the oxidizer and High-Density PolyEthy  lene as fuel; the configurations included free-jet test, in which small button-like samples were exposed   to the supersonic exhaust jet of the rocket nozzle; and chamber inserts, in the shape and size of an an  nular element, placed inside the rocket combustion chamber. Computational Fluid Dynamic simulations,   for modeling heat transfer and combustion chemical reactions, complemented the experimental observa  tions and supported the characterization of test conditions. Samples with ZrB 2 -SiC matrix and continuous   or chopped carbon fibers, sintered by either Hot Pressing or Spark Plasma Sintering were tested. Free  jet test samples demonstrated a substantially improved erosion resistance with respect to conventional   graphite and in one case a negligible material recession. UHTCMC samples erosion was associated to the   occurrence of a rapid rise in surface temperature, which achieved values over 2900 K. Chamber inserts,   besides confirming the outstanding erosion resistance of UHTCMCs with respect to traditional materials   (i.e. C/SiC), proved that long-fibers samples with sufficient porosity are more likely to withstand thermal   shocks typical of the rocket combustion environment.   © 2020 Elsevier Ltd. All rights reserved.   1. Introduction   fects over the motor operation. Thus, the requirement that dimen  sional stability of the nozzle throat should be maintained makes   Challenges for solid and hybrid rocket technologies include the   the selection of suitable rocket nozzle materials extremely hard.   optimization of thermal insulations and the design and fabrica  The materials used for these applications include refractory   tion of non-eroding firing thrusters able to survive severe thermal  metals, refractory metal carbides, graphite, ceramics, carbon  structural and thermal-chemical combustion environments without   carbon composites and fiber-reinforced plastics [1-3] . Certain   cooling systems. The inner surface of the exhaust nozzle, through   classes of materials demonstrated superior performances under   which the propellant flow is accelerated to supersonic conditions,   specific operating conditions, but the choice depends on the spe  is very critical in this sense, as it is subjected to the highest shear   cific application. For instance, fully densified refractory-metal noz  stresses, pressure and heat fluxes in a chemically aggressive envi  zles generally are more resistant to erosion and thermal-stress   ronment. These severe conditions usually lead to removal of sur  cracking than the other materials. Graphite performs well with   face material (ablation) due to heterogeneous reactions between   the least oxidizing propellant but is generally severely eroded and   oxidizing species in the hot gas and the solid wall. Because of the   must be replaced after each run [4-6] .   material erosion, enlargement of the nozzle throat section occurs   In recent years, Ultra-High-Temperature Ceramic (UHTC) mate  with consequent decrease of the rocket thrust and detrimental ef  rials, including zirconium or hafnium diborides or carbides, have   ∗  Corresponding author.   E-mail address: stefano.mungiguerra@unina.it (S. Mungiguerra).   https://doi.org/10.1016/j.ijheatmasstransfer.2020.120492   0017-9310/© 2020 Elsevier Ltd. All rights reserved.   assumed an increasing importance in aerospace research because   of their high temperature capabilities, with melting points above   3500 K, high temperature strength and oxidation resistance at   service temperatures exceeding 2300 K. Some of these materials   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   proved to be very interesting to develop aerospace components   working in harsh environments [7-9] . The use of single-phase ma  terials, without secondary phases, is not sufficient because they are   characterized by low fracture toughness, low thermal shock resis  tance and lack of damage tolerance [10] . To improve their prop  erties, UHTC composites with SiC or other Silicon based ceram  ics, in the form of particles, short fibers or whiskers have been   developed with better tolerance and thermal shock resistance in   aggressive chemical environments [11-13] . Unfortunately, despite   the very good oxidation resistance of small specimens, larger UHTC   components frequently exhibit poor reliability and were subject to   catastrophic failures in high enthalpy flows. Based on these re  sults, the current research activities are oriented towards Ultra  High-Temperature Ceramic Matrix Composite (UHTCMC) materials   containing high volume fraction of C or SiC fibers in UHTC matri  ces, which represent the next step to bring significant technolog  ical improvements in comparison to the state-of-art materials for   aerospace applications [14,15] .   In this framework, University of Naples “Federico II”  (UNINA)   and the Institute of Science and Technology for Ceramics (ISTEC)   are involved in the Horizon 2020 European C  3 HARME research   project, focused on a new class of UHTCMCs for near zero-erosion   rocket nozzles. In particular, an extensive experimental character  ization campaign is ongoing, based on an incremental approach,   envisaging tests on prototypes of increasing complexity, from small   button-like samples to complete sub-scaled De Laval nozzles [16] .   Several activities have been already carried out for the char  acterization of mechanical properties at room and high tempera  ture. Specifically, flexural strength values as high as 450 MPa and   20 0-30 0 MPa were collected at 1500 °C and 2100 °C, respectively,   demonstrating the excellent performance of UHTCMCs based on   Carbon fibers preforms and HfC, ZrC, TaC and ZrB 2   matrices [17-  19] . All the high temperature values were always higher than the   room temperature values. Moreover, UHTCMCs based on carbon   preforms and ZrB 2   matrix displayed also excellent thermal shock   resistance [19] , with 85% of the pristine flexural strength retained   after water quenching from 1500 °C to room temperature and even   remarkable strength retention upon exposure to oxyacetylene torch   at 2240 °C [20] .   To improve the achieved Technology Readiness Level and pro  vide a more comprehensive material characterization, the experi  mental activities presented in this work are aimed at screening the   most suitable material candidates for the production of the final   large nozzle, by evaluation of oxidative, mechanical and ablation   resistance in a relevant environment for the final application. Two   novel, dedicated test configurations were therefore developed and   successfully employed at UNINA Aerospace Propulsion Laboratory   to evaluate the potentialities of different ZrB 2  -based UHTCMC sam  ples manufactured by ISTEC. A 200 N-class hybrid rocket was used,   selecting as propellants gaseous oxygen as oxidizer and a cylindri  cal grain of solid High-Density PolyEthylene (HDPE) as fuel, repro  ducing a typical hybrid rocket combustion environment.   In the first set-up (“Free-jet test”), small button-like samples   were exposed to the supersonic exhaust jet of the rocket nozzle.   The main objective was to assess the capability of the material to   withstand harsh aero-thermo-chemical environment in moderate   aerodynamic conditions, in a stagnation point configuration: heat   flux in the order of 15 MW/m  2 , temperature above 2500 K, pres  sure in the order of 3-4 bar, oxidizing and chemically reacting at  mosphere.   The second set-up (“Chamber insert test”) was meant to as  sess the capability of the test article (in the shape and size of an   annular combustion chamber element) to withstand high thermo  mechanical stress at high pressures (order of 10 bar) in relevant   aerothermo-chemical combustion environment (heat flux around 5   MW/m   2 ).   The combined effort of experimental diagnostics and Compu  tational Fluid Dynamic (CFD) simulations, able to predict the flow   field around the test articles and the corresponding aero-thermo  dynamic loads, provided an insightful characterization of the op  erating conditions and an improved understanding of the experi  mental observations. Material performance was evaluated both in   terms of erosion resistance and of post-test structural integrity.   Post-test inspections were also carried out to analyze the sur  face modifications occurred after the exposure to the aero-thermo  chemically aggressive combusting flow.   2. Materials and methods   2.1. Experimental setup   The test rig is a versatile set up primarily designed for testing   hybrid rocket engines of several sizes [21] . It is equipped with a   test bench and a general-purpose acquisition system, which allows   evaluating propellant performance and combustion stability [22] ,   testing of sub-components and/or complete power systems, noz  zles [23] , air intakes, catalytic devices [24] , burners, ignition and   cooling systems [ 10 , 25 ].   The schematics of the rocket employed in this work is depicted   in Fig. 1 , which also shows the UHTCMC sample for free-jet test   behind the rocket nozzle, and the combustion chamber annular   disk insert downstream the fuel grain.   A detailed description of the laboratory and of the experimental   facilities can be found in Ref. [26] . The tests presented in the fol  lowing sections were performed with a converging nozzle injector,   whose exit-section diameter is 6 mm, which delivered oxygen in   single-port cylindrical fuel grains of HDPE. A graphite converging  diverging exhaust nozzle with a throat diameter of 10.4 mm and   an area ratio of 2.54, is employed to expand the combustion prod  ucts up to supersonic speeds.   Oxidizer mass flow rate is measured by means of a Tescom   ER30 0 0 pressure regulator, which controls an electropneumatic   valve in order to reduce the pressure to the desired setpoint, up  stream a chocked Venturi nozzle. The average fuel mass flow rate   is estimated by means of grain mass measurements before and af  ter the test (see Ref. [27] ).   The two test set-ups developed for the current research activi  ties will be described in the following subsections.   2.1.1. Free-jet test   The Free-jet test configuration was conceived to expose differ  ent UHTCMC samples to the supersonic exhaust jet of a 200 N  class hybrid rocket nozzle. The specimens were small, button-like   samples, with maximum diameter of 17 mm ( Fig. 2 ).   Each specimen was placed downstream the hybrid rocket en  gine, in order to be reached by the exhaust plume coming from   the nozzle (see Fig. 1 ). The size of the specimen, chosen accord  ingly to the available mechanical interfaces, is such that the front   surface diameter is comparable to the nozzle exit section, allowing   the sample to be thoroughly hit by the supersonic jet. The exper  imental set-up consisted in a mechanical system connected to the   test bench in order to support and keep the specimen aligned with   the motor axis. This system was designed to place the test article   at the desired distance to the nozzle exit, d . In the present test   campaign, a distance of 15 cm was selected, compatibly with the   mechanical requirements for the fixation to the test bench, in or  der to provide realistic conditions to test the materials, in terms of   heat fluxes, temperatures, pressures and gas composition. In fact,   it is expected that, for this kind of applications, throat heat fluxes   in hybrid rockets are on the order of 15 MW/m  2 , throat pressures   are on the order of a few bars, temperatures can be higher than   30 0 0 K [ 16 , 28 , 29 ], while the reacting gas products are the same, as   2   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 1. Layout of 200 N-class hybrid rocket engine, including the set-up for UHTCMC testing.   emittance versus testing time. On one hand, the two-color pyrom  eters overcome this problem measuring the true temperature. On   the other hand, the IR-TC detects the spectral radiance coming out   from the heated sample along the infrared band wavelengths of  0.8-1.1  μm. The surface temperature distribution can be calculated   assuming constant emissivity along the monitored surfaces of the   samples and taking into account the axial symmetry of the spec  imens. Once the local temperature is measured thanks to the py  rometer at the measurement spot, that value is input to determine   the spectral emittance in the range of the IR-TC, and finally the   surface temperature distribution is evaluated.   2.1.2. Combustion chamber insert test   The second UHTCMC sample was designed as a simple flat disk.   Fig. 4 shows the prototype design, characterized by a simple shape:   an external diameter of 68 mm, an internal diameter of 15 mm and   a thickness of 5 mm. Referring to Fig. 1 , the new-class material   insert can replace one of the thermal protections usually placed   between the fuel grain and the post-chamber.   This chamber insert was conceived to combine simple shape   and low manufacturing cost, allowing to evaluate the response of   the materials to be tested in typical operating conditions in the   combustion chamber of a hybrid rocket, where high temperatures   and pressures are expected, in presence of a combustion flame,   with relevant chemical composition. Although a relevant environ  ment for nozzle applications was not fully reproduced (gas veloci  ties are low in the combustion chamber), this kind of sample can   be considered a breadboard validating the main structural func  tionalities of the final component in a laboratory-reproduced rep  resentative environment.   2.2. Material processing   Fig. 2. Nominal design of UHTCMC samples for free jet tests. Dimensions are in   mm.   in free-jet test the flow comes directly from nozzle exit. These con  ditions are well reproduced by the present configuration, as will   be shown in detail in Section 2.4 . The described configuration was   in some way similar to that adopted by the authors to test ma  terials in an arc-jet wind tunnel simulating the re-entry conditions   [30,31] , simplifying samples manufacturing; but while in this latter   case the environment is characterized by near-vacuum conditions   (stagnation pressure on the sample of the order of 10 kPa), in the   free-jet tests presented in this work the stagnation pressure was   around 4 bar. Therefore, the material was subjected also to signifi  cant structural stresses due to the relatively high pressure.   Fig. 3 shows pictures of the experimental set-up, including the   non-intrusive diagnostic equipment employed for the real-time   evaluation of the sample surface temperature. In particular, the   surface temperature of the samples was continuously measured  ±1% instrumental accuracy) by digital two-color pyrometers (In (   fratherm ISQ5 and IGAR6, Impac Electronic GmbH, Germany) at an   acquisition rate of 100 Hz. In addition, the infrared response of   Commercial products were used to prepare the UHTCMC, where   the specimen during the free-jet testing was obtained by means   of an infrared (IR) thermo-camera (TC, Pyroview 512 N, DIAS In  frared GmbH, Germany). The two-color ISQ5 pyrometer exploits  two overlapping infrared wavelength bands at 0.7-1.15  μm and  μm to measure the temperature from 1273 K up to   0.97-1.15   3273 K. The IGAR6 pyrometer operates in the bands 1.5-1.6 μm   and 2.0-2.5 μm to return the sample temperature in the range   the matrix was ZrB 2  (H.C. Starck, Germany, Grade B, d 50  : 2.3  μm, impurities, wt%: O:0.8, Hf:1.82, N:0.19) added with 10 vol%  μm, impurities, wt%:   SiC (H.C. Starck, Germany, UF-25, d 50  : 0.42   O:1.7, Fe:0.04, Al:0.03). Pitch-derived carbon fibers were used: ei  ther Unidirectional (UD) fabrics (Nippon Graphite Fiber, Japan, UF  XN80-300) or chopped (Nippon Graphite Fiber, Japan, XN-80C-03S  with diameter of 10  μm and chopped length of 3 mm)   523-2273 K. The two pyrometers, both pointing at the central area   Materials containing long fibers were processed as described in   of the samples front surface, gave perfectly equivalent responses,   [14] , by following slurry infiltration of carbon fiber fabrics, stacking   so only the temperature profiles measured by ISQ5 are herein re  in 0/90 ° configuration, drying and sintering. As for short fiber com  ported. The pyrometers mode can be set in order to give back the   posites, chopped fibers were milled to the ceramic powders by fol  peak value of the temperature field detected inside the measure  lowing conventional wet ceramic procedures including ball milling,   ment area, consisting in a round spot of roughly 5 mm in diam  drying in rotary evaporator and sintering in either hot pressing   eter. In addition, the so-called “two-color mode” provides an out  (HP) or spark plasma sintering (SPS), as reported in Table 1 and   put value independent on the (directional) spectral emittance. It   2 [11] .   is generally assumed that the observed surface behaves as a gray   Three samples were manufactured for free-jet tests and two   body over the operating temperature range. Surface chemical re  chamber inserts were manufactured, with their main characteris  actions occurring during test can be responsible for changes in   tics summarized in Table 1 and Table 2 , respectively.   3   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 3. Set-up for free-jet test. The area within the red circle in the left picture is zoomed in the right picture. (For interpretation of the references to colour in this figure   legend, the reader is referred to the web version of this article.)   Table 1   Free-jet test samples.   Label   Matrix   Carbon fibers   Sintering   Porosity   type   vol%   m ethod   °C/MPa/min   vol%   FJ-LF-1   ZrB 2   + + +   10 vol% SiC   XN80 0/90 °  45   SPS   1900, 40, 5   <   1   FJ-LF-2   ZrB 2    10 vol% SiC   XN80 0/90 °  45   HP   1900, 40, 10   11   FJ-SF   ZrB 2    10 vol% SiC   XN80 chopped   45   SPS   1900, 40, 5   <   1   Table 2   Chamber insert samples.   Label   Matrix   Carbon fibers   Sintering   Porosity   type   vol%   m ethod   °C/MPa/min   vol%   CI-LF   ZrB 2   + +   10 vol% SiC   XN80 0/90 °  45   HP   1900, 20, 15   15   CI-SF   ZrB 2    10 vol% SiC   XN80 chopped   45   SPS   1900, 40, 5   <   1   Fig. 4. Nominal design of UHTCMC combustion chamber inserts.   2.3. Numerical models   2.3.1. One-dimensional model for combustion chamber conditions   simulation   A one-dimensional model based on NASA CEA software [32] can   be used to rapidly evaluate the evolution of the operating condi  tions in the combustion chamber, in particular the chamber pres  sure, and through the nozzle during the time. In this case, the in  put data of the model are the oxidizer mass flow rate, the geomet  rical dimensions of the fuel grain and the operating time.   As the instantaneous regression rate is an unknown parameter   and the oxidizer mass flux and chamber pressure depend on the   regression rate itself, the expected data were estimated assuming   the classical regression rate law   ˙ r   =   aG   n  ox   (1)   where the coefficients a and n were selected from the values avail  able in literature relevant to the combustion of gaseous oxygen   with HDPE fuel grains [33] . Integrating Eq. (1) in time, the instan  taneous port diameter D ( t ) was calculated. Then, considering the   prescribed oxidizer mass flow rate, the corresponding mass flux   G ox ( t ) and regression rate ˙ r   (t )   was estimated. Then, the fuel mass   flow rate was calculated as   ˙ m f   (   t )  =  ρ  f   π D (   t )  L ˙ r (  t )  (2)   where   ρ  f   is the solid fuel density and L is the length of the grain,   and correspondingly the average mixture ratio OF   (t )  =  ˙ m ox  ˙ m f   (t )  can   be derived. From these calculations, the estimation of the aft  chamber pressure p c was performed by means of an iterative pro  cedure to solve the steady-state mass balance equation   ˙ m ox   A t   (cid:2)  1   +  1   OF   (cid:3)  =  p c   ηC   ∗  (3)   in which A t is the nozzle throat area, C   ∗  is the theoretical char  acteristic exhaust velocity (that primarily depends on the mixture   ratio and, to a minor degree, on pressure) and the combustion ef η , assumed equal to unity. For the dependence of the C   ficiency,   ∗  4                   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Table 3   C 2 H 4 - O 2 reaction system.   No.   Reaction   a   A K  b   β  E a  b   1   C 2 H 4   +   O 2   (cid:2) 2CO  (cid:2) CO 2  (cid:2) CO 2   OH  (cid:2) OH  (cid:2) OH   +   2H 2   1.80e   + + + + + + + + + +   14   0.0   −4540  35,500  −740   2   CO   + + +   O   +   M   5.30e    13   0.0   3   CO   + + +   M   4.40e    06   1.5   4   H 2    O 2    OH   1.70e    13   0.0   48,000   5   H + O 2    O   2.60e    14   0.0   16,800   6   OH   +   H 2   (cid:2) H 2 O   +   H   2.20e    13   0.0   5150   7   O + H 2   (cid:2) OH   +   H   1.80e    10   1.0   8900   8   OH   +   OH   (cid:2) H 2 O  (cid:2) H 2  (cid:2) H 2 O   +   O   6.30e    13   −1.0  0.0  −2.0   1090   9   H + H   +   M   6.40e    17   0   10   H   +   OH   +   M   2.2e    22   0   a Third-body efficiencies for all termolecular reactions are 2.5 for   M   =   H 2 , 16.0 for M   H 2 O, and 1.0 for all other M.  b Units are in seconds, moles, cubic centimeters, calories and Kelvin.   =  on pressure, Eq. (3) is implicit and an iterative calculation tech  nique is needed. A combustion pressure was first assumed, then   the CEA code was run to calculate the equilibrium composition and   the theoretical exhaust velocity, assuming frozen flow through the   nozzle, at the given OF ratio in input. Finally, combustion pressure   was adjusted repeatedly until convergence.   2.3.2. CFD model for the simulation of flow and heat transfer in   free-jet test   In order to provide an improved understanding of test condi  tions, CFD simulation of the external plume of the exhaust gases   was performed. The model herein presented has been already suc  cessfully employed and validated [28] for a preliminary free-jet   test on a graphite sample used as a reference in this work (see   Section 3.1 ). A further validation of the model can be found in the   Supplementary material.   As inlet boundary conditions, the time-averaged results of the   numerical tool described in the previous section were employed.   Specifically, the average conditions achieved at nozzle exit section   were used as input for the simulation of the flow field downstream   the rocket.   To this purpose the Reynolds-Averaged Navier-Stokes (RANS)   equations for single-phase multicomponent turbulent reacting   flows were solved with a control-volume-based technique and a   density-based algorithm [34] , employing the Shear Stress Transport  (SST) k -ω  model as turbulence closure [35] . A detailed analysis of   thermo-chemical evolution of gas mixture was performed, in order   to have an accurate prediction of heat transfer at solid walls. The   transport equations for the main combustion products (O 2   , C 2  H 4   ,   H 2   O, CO 2  , CO, H 2  , H, O, OH are the species considered in the cur  rent model, together with the non-reacting N 2  ) were solved, and   the Eddy Dissipation Concept (EDC) model [36] was employed for   the combustion mechanism, which accounts for detailed chemical   reaction rates in turbulent flows. Consequently, the Arrhenius rate   K for each reaction was calculated as   K   =   A K T   β  exp   (cid:2)  − E a   RT   (cid:3)  (4)   where A K is a pre-exponential factor, T is temperature,  β is a tem  perature exponent, which allows taking into account reaction rate   dependence on temperature, E a is the activation energy and R is   the universal gas constant. The constants were taken from Ref.   [37] and are reported in Table 3 .   The Discrete Ordinates model [38] for the radiation is included   in the numerical modeling.   The computational grid used for the simulation of the free re  acting jet exiting from the nozzle is shown in Fig. 5 . A pressure   inlet boundary condition was set on the surface representative of   the nozzle exit section, imposing the total pressure and the total   temperature corresponding to the operating chamber pressure and   temperature in the rocket, and the static pressure and the chem  ical composition at the exit of the nozzle. The ambient pressure   was set on the other external boundaries of the computational do  main.   2.3.3. CFD model for the simulation of flow and heat transfer in   chamber insert test   The application of suitable numerical model assumes even a   higher importance for having a deep insight on the operating con  ditions for the test of chamber inserts, which are located inside the   combustion chamber of the rocket and therefore cannot be moni  tored by means of optical diagnostic instruments used in the case   of free-jet tests.   Also in this case CFD simulation were carried out applying the   numerical model specifically defined by the authors at UNINA for   the detailed simulation of hybrid rockets internal ballistics. This   has been extensively validated in several works based on a sound   comparison between numerical and experimental results obtained   by means of the same engine of the present paper and several pro  pellants (including GOx-HDPE) [29,39-41] . In those works, the cor  rect estimation of the wall heat fluxes, which is the main objective   of the present simulations, is a key factor to correctly predict the   fuel regression rate, and the achieved agreement between simula  tion and experiment has been highly satisfactory in a range of test   cases.   Again, the model provides the resolution of the RANS equa tion considering the SST k -ω  model as turbulence closure. In this   case, considering that the diffusion processes typically occurring in   hybrid rocket motors were slow with respect to the chemical ki  netics, the non-premixed combustion of oxygen and gaseous fuel   injected from the wall was modeled by means of the Probability   Density Function (PDF) approach coupled to chemical equilibrium   [42] . Accordingly, combustion was simplified to a mixing problem,   and the difficulties associated with closing non-linear mean re  action rates were avoided. The turbulence-chemistry interactions   were described by means of the average mixture fraction, f , and   its variance, f’  2 . The shape of the assumed PDF was described by  β -function of these two quantities. Once f and f’  2 were cal  the   culated at each point in the flowfield, the known PDF was used   to compute the time-averaged values of individual species mole   fractions, density, and temperature with simple thermochemistry   calculations based on the minimization of Gibbs free energy [32] .   Although the actual products of solid fuel pyrolysis are numerous   and their composition depends on both the wall temperature and   heating rate, the practice common in the literature is to consider   the gaseous monomer alone; here, thus, gaseous ethylene were   considered as the fuel injected from the grain surface.   Fig. 6 shows a typical axisymmetric computational grid em  ployed for the CFD simulation of the flow field inside the sub  scale hybrid rocket, constituted by the internal volume of the pre  chamber, of the fuel grain port, of the aft-mixing chamber and of   the nozzle.   The model includes a specific treatment for the boundary con  ditions at the interface between the gaseous flow region and the   solid fuel wall, based on the mass, energy and species balance and   an Arrhenius-type equation for the fuel pyrolysis, in order to allow   calculation of the local fuel regression rate, which governs the fuel   mass flow rate and in the last instance the oxidizer-to-fuel ratio.   2.4. Test conditions   This section presents the nominal test conditions for the two   experimental configurations. All tests had a nominal duration of   10 s. Cylindrical 220 mm long HDPE grains were employed as fuel   and gaseous oxygen as oxidizer. Initial fuel grain port diameter was   15 mm. For free-jet tests, a nominal oxidizer mass flow rate of   5   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 5. Computational grid for the simulation of the free reacting jet exiting from the rocket nozzle.   Fig. 6. Typical computational grid for the simulation of hybrid rocket internal ballistic.   Table 4   Table 6   Summary of nominal free-jet (FJ) and chamber insert (CI) test conditions.   Summary of conditions at sample location.   Test condition FJ   Test condition CI   Stagnation point pressure [bar]   Oxidizer mass flow rate [g/s]   Average Oxidizer-to-Fuel ratio   Chamber pressure [bar]   30   5.63   6.43   40   6.50   8.2   Combustion temperature [K]   ~ 3200   ~ 3200   Average CO 2 mass fraction   Average H 2 O mass fraction   Average O 2 mass fraction   Average Cold-wall surface heat flux [MW/m   2 ]   Peak Cold-wall surface heat flux [MW/m   2 ]   3.8   0.29   0.12   0.31   12.9   16.0   Table 5   Aero-thermo-chemical conditions   at nozzle exit for free-jet tests.   Pressure [bar]   0.55   Temperature [K]   ~ 2300   Mach Number   CO 2 mass fraction   H 2 O mass fraction   O 2 mass fraction   2.22   0.35   0.16   0.34   In particular, Fig. 7 (a) confirmed that the temperature was higher   than 30 0 0 K, which is close to the value achieved in the com  bustion chamber, evaluated by means of the chemical equilibrium   software. Fig. 7 (b) shows instead that, downstream a sequence of   shock/expansion oblique waves, the stagnation-point peak pres  sure was almost 4 bar. Finally, Fig. 7 (c) shows the distribution of   O 2   mass fraction, whose average value at sample location is 0.31,   demonstrating that there was still a significant amount of oxidizer   in the flow, causing a consistent material oxidation, as detailed be  30 g/s was used, whereas for chamber insert test the oxidizer mass   low.   flow rate was set to 40 g/s. Table 4 summarizes the main operating   The thermo-fluid-dynamic and chemical flow field produces, on   parameters in the rocket combustion chamber for the abovemen  the sample front surface, the cold-wall heat flux, static pressure   tioned test conditions, calculated by means of the one dimensional   profile and oxygen mass fraction shown in Fig. 8 . The peak heat   model presented in Section 2.3.1 .   flux is 16 MW/m  2 , which is a value representative of the actual   Based on the results of 1D simulations, the model described   loads that might be expected on a hybrid rocket nozzle throat sur  in Section 2.3.2 was employed to predict, by CFD simulations,   face [16] . Moreover, it is possible to see that heat flux, pressures   the flow field in the exhaust plume of the nozzle, for the free  and oxygen concentration are maximum on the symmetry axis. As   jet test set-up and the corresponding loads on the material sam  it will be extensively discussed in Section 3.1 , this is the location   ple. The thermo-fluid-dynamic and chemical conditions at nozzle   where a temperature jump phenomenon can be triggered, associ  exit section, evaluated by means of the 1D model, are reported   ated to a consistent local sample erosion.   in Table 5 and were set as boundary condition for the inlet sec  Table 6 summarizes the main thermo-fluid-dynamic and chem  tion of the CFD domain. Fig. 7 shows the distributions of tempera  ical conditions at sample location, including the amount of other   ture, pressure and molecular oxygen mass fraction in the flow field.   6   oxidizing species (CO 2   , H 2   O) during free-jet tests.   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 7. (a) Static temperature, (b) static pressure, (c) O 2 mass fraction distributions around the sample in free-jet test.   Fig. 9. CFD results around the chamber insert: temperature contour plot with over  lapped streamlines (top half) and mixture-fraction isolines (bottom half).   Table 7   Average operating conditions in the rocket   post-chamber.   Total Pressure [bar]   Average Oxidizer-to-Fuel ratio   7.90   6.15   Average Total Temperature [K]   2900   Fig. 8. Cold-wall total surface heat flux, static pressure and O 2 mass fraction on   free-jet sample front surface, calculated by CFD.   mass-fraction in the unburned mixture isolines drawn on the bot  tom half, in the region surrounding the chamber insert, located, as   mentioned before, between the fuel grain and the post-chamber.   From this picture, all the main features of the internal flowfield   can be unveiled. For example, it should be noted that the inclusion   Similarly, the CFD model for the simulation of hybrid rocket in  of the chamber insert and the corresponding geometrical discon  ternal ballistics described in Section 2.3.3 was employed to esti  tinuity determined a slight change in the fluid behavior, which, as   mate the operating conditions to which the chamber inserts were   also pointed out in the literature [43] , further promoted the pro  subjected during the test.   pellant mixing and the combustion efficiency in the post-chamber.   Fig. 9 shows the plot of the calculated temperature contour   Table 7 summarizes the average operating conditions estimated   with the streamlines overlapped on the top half, and the fuel   in the transversal section corresponding to the chamber insert lo  7   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 10. Thermal histories of the free-jet test samples, recorded by the pyrometer.   Fig. 12. Erosion rates of free-jet test samples.   Table 8   front surface area, obtaining the equivalent axial average erosion   Summary of conditions on the chamber insert surface.   rate.   Average CO 2 mass fraction   Average H 2 O mass fraction   Average Cold-wall surface heat flux [MW/m   2 ]   Peak Cold-wall surface heat flux [MW/m   2 ]   0.58   0.25   5.8   7.8   3. Results and discussion   3.1. Free-jet tests   cation, which are in good agreement with the values obtained with   whereas the time-averaged combustion chamber pressure was   the one-dimensional model (see Table 4 ). Finally, Table 8 reports   about 6.5 bar. Fig. 10 shows the time history of the maximum tem  the conditions on the sample surface, on which a cold-wall average   heat flux equal to 5.8 MW/m  2 is experienced, with a peak value of   7.8 MW/m  2 . Since in the typical chemical field of hybrid rockets   perature reached by the samples, detected by the ISQ5 pyrome  ter, operated in two-color mode. Although the test conditions were   the same, the final temperature reached by the sample FJ-LF-2 was   the oxidizer is mainly concentrated in the core flow, the oxidizing   over 400 K lower than that reached by the other two samples. In   species interacting with the insert surface were prevalently com  fact, after 4-5 s, the temperature of FJ-LF-1 and FJ-SF started ris  Measured oxidizer mass flow rate was 31 g/s in all the test,   bustion products like CO 2   and H 2  O [44] rather than O 2  .   2.5. Materials performance assessment   ing more rapidly, reaching in the end values over 2800 or even   2900 K. In the final part of the test, after the sudden temperature   jump, solid fragments were wiped off  the samples surfaces by the   oncoming supersonic flow, as visible in Fig. 11 , showing, as an ex  ample, pictures recorded by the optical camera at the beginning   Materials capability to withstand the harsh conditions above   and at the end of test on sample FJ-LF-1. A similar phenomenon   described was first evaluated by structural inspection and then by   was observed for the short-fibers sample, whereas no instability   microstructure analysis on the surfaces and polished cross-sections   was detected in the case of FJ-LF-2. Fig. 12 shows the samples   by field emission scanning electron microscopy (FE-SEM, Carl Zeiss   equivalent average erosion rates, evaluated by means of mass mea  Sigma NTS GmbH, Oberkochen, DE) and energy dispersive x-ray   surements before and after test. Data are also compared to that   spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK).   related to the test of a graphite sample, used as a reference mate  Free-jet test samples mass was measured before and after test, and   rial, whose details are reported in [28] . It is clear that the samples   an equivalent average erosion rate, expressed in mm/s, was calcu  that reached the ultra-high temperatures were subjected to consid  lated based on mass loss. The latter was converted in a volume loss   erable erosion, whereas sample FJ-LF-2 preserved stable mass and   dividing it by material density, then divided by actual test duration.   dimensions. It needs however to be remarked that the erosion rate   Finally, assuming a uniform consumption of the sample in the ax  of the two UHTCMC samples that underwent erosion, FJ-LF-1 and   ial direction, the volume loss rate was divided by the specimen   FJ-SF, was over 2 times lower than that of graphite, increasing the   Fig. 11. Pictures of test on FJ-LF-1 sample, before (left) and after (right) the temperature jump.   8   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 13. Pictures of samples (a) FJ-LF-1, (b) FJ-LF-2 and (c) FJ-SF before (top) and after (bottom) the test.   confidence in the fact that an optimized formulation of ZrB 2  -SiC  The result is presented in Fig. 14 , where it is shown that in the   based materials could be a significant improvement with respect   earliest stages of the test, the emissivity had a steep increase to   to state-of-art materials.   values up to 0.9, then slightly decreased to 0.8, and finally, when   Fig. 13 shows pictures of the three samples before (top) and   the jump was triggered, it appeared to increase again, keeping, in   after (bottom) the test. All of them appeared oxidized after ex  posure to the supersonic flame, but, whereas the oxide layer of   the final part of the test, vales close to a black body radiator. Ap ελ , some images were   plying the corresponding correct values of   FJ-LF-2 survived the thermo-mechanical load, preserving structural   extracted from the IR-TC. Specifically, Fig. 15 (a) shows the tem  integrity, the FJ-LF-1 and FJ-SF underwent notable shear stresses   perature distribution over the sample surface at different time in  and their shape was consistently modified upon achieving temper  stants, around the moment in which the abrupt temperature rise   ature peaks of 2860 and 2930 K, respectively. A valley formed in   occurred, and Fig. 15 (b) reports the corresponding temperature ra  the FJ-LF-1, Fig. 13 a, whilst the FJ-SF ended with a lateral squash  dial profiles. The figure highlights that the temperature started ris  ing, Fig. 13 c, possibly due to a slight misalignment that resulted   ing suddenly in a localized region in the center of the surface, pro  fatal under the high thermo-mechanical loads.   gressively spreading its area. As also highlighted above by numer  A detailed analysis of the thermographic images was carried out   ical simulations, this is the area where the heat flux was most in  for sample FJ-LF-1, to better understand the mechanism observed   tense and the concentration of molecular oxygen was expected to   during the jump, which led to the high erosion rate. First, as ex  be higher, so the jump might be associated to triggering, at high   plained in Section 2.1.1 , the sample spectral emissivity had to be   temperature, of a consistent sample oxidation; the porosity left by   estimated in order to get reliable quantitative data from the IR-TC.   depletion of carbon fibers and the characteristic low thermal con  9   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   tive properties might have got closer to those of a perfect black   body.   3.1.1. Post-test microstructure analysis   The surface of the FJ-LF-1 sample was very rough and whitish,   no carbon fiber was left in any position of the model that was   completely covered with ZrO 2   which displayed diverse morphol  ogy, either dense and coarse, with brighter contrast in Fig. 16 a, or   finer, resulting in a more gray contrast, Fig. 16 b,c. In the outer re  gions, some isolated SiO 2   droplets could be found, Fig. 16 c. The   central area, which remained depressed by the jet, displayed some   round pits attributable to the exposure and consumption of fiber   bundles, Fig. 16 d. In these pits, ZrO 2   was very fine and covered   with B 2   O 3   crystals, as a result of the rapid oxidation and cooling   of the boride matrix, Fig. 16 e. Notably, several melts could have   formed on the surface which touched the peak temperature of   2860 K. Silicon oxide exhibits a liquid phase already above 1690 K   [47] and from this temperature the oxidized surface of the ZrB 2   matrix, oxidized into ZrO 2  , started to dissolve into SiO 2  glass form  ing a Si-O-Zr melt. With further heating, the solubility of ZrO 2   in  Fig. 14. Near-infrared spectral emissivity of sample FJ-LF-1 during the heating   phase of the test, compared with the temperature measured by the ISQ5 pyrom  eter.   ductivity of oxidized phases such as ZrO 2   [ 45 , 46 ] (further details   creased gradually. According to the ZrO 2  -SiO 2   equilibrium phase   on microstructure and oxidation will be given in next section), to  diagram [48] , the melting temperature of zirconium oxide, 2988 K,   gether with the peaked distribution of heat flux along the surface,   can be reduced by presence of liquid silicon oxide, by meeting an   might have favored the steep increase in temperature in the cen  tral area, therefore fostering considerable thermo-mechanical ero  eutectic temperature already at 1960 K. While ZrO 2   reaches high   solubility in SiO 2  at 2520 K, SiO 2  becomes unstable when the tem  sion. The corresponding increase in the spectral emissivity could   perature overpasses 2134 K and starts to leave the system in form   be therefore associated to the formation of a cavity, whose radia  of SiO gas. The progressive evaporation of SiO 2  makes the ZrO 2  sol  Fig. 15. (a) IR thermal images of sample FJ-LF-1 at different time instants, around the moment in which the abrupt temperature rise occurred, and (b) corresponding   temperature radial profiles.   10   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 16. SEM images of the surface of the FJ-LF-1 sample after test showing a) the overall aspect with different sized ZrO 2 grain zones, b) melted/coarse or c) fine with SiO 2   scattered droplets and d) magnification of the pits with enlarged views in e) and f). EDS of ZrO 2 and B 2 O 3 are reported below. (For interpretation of the references to colour   in this figure legend, the reader is referred to the web version of this article.)   ubility decrease and its precipitation as fine grains is favoured. The   face was whitish, but rather smooth as compared to the FJ  following coalescence and coarsening is a natural process occurring   LF-1. The external surface displayed dark-contrasted zones, fea  at such extreme temperatures.   tured by the presence of silica coverage where ZrO 2   precipi  Looking at the cross section in Fig. 17 a, the valley can be bet  tated and bubbles formed owing to outgassing from the subscales,   ter appreciated and its formation can be associated to migration   of material to the edges which lost their sharp shape, Fig. 17 b.   Analysis of the core of the button evidenced partial detachment of  μm thick, Fig. 17 c&g,   an outermost dense oxide scale, around 50   Fig. 19 a&b, like typically observed for conventional ZrB 2  -SiC ox  idized materials [51] . Other zones were instead richer in ZrO 2   with scattered glass islands, Fig. 19 a&c. To note that these ar  eas, where an elongated fibrous morphology is recognized, were   but this scale left an almost unaffected boride matrix, where only   allegedly zones where the UHTC-impregnated carbon fiber fabric   fiber/matrix interaction became more intimate, Fig. 17 c,f&h. The   was exposed after machining of the model, owing to slight mis  fiber remained in their site, but their edge turned more irregular   alignment of the fabric layers during shaping and hot pressing.   and jagged, similar to a situation where over-sintering took place,   When high fraction carbon-based material was exposed to the di  resulting in a very strong interface and compromised fiber struc ture, like observed in [49] . About 170  μm underwent this modifi  rect jet impingement, partial oxide detachment and chipping was   observed, Fig. 19 c.   cation in the central zone. The jet had indeed the effect of a fur  The good response of this material to the jet flow was con  ther sintering in the sub-scale layers. In addition, the extremely   firmed after analysis of the cross section, whose overview is dis  high temperature achieved, 2860 K, was well above some eutec  played in Fig. 20 a. The flat shape was maintained with only little   tic melting temperatures found in the quasi binary ZrB 2  -SiC and  ZrB 2  -C tie lines at 2400 and 2673 K, respectively [50] . As confir  puckering of the profile, Fig. 20 b. The outermost oxide scale, less  than 120  μm, was composed of dense ZrO 2   with tiny channels of   mation of the passage through these eutectics, SiC phase was pre  SiO 2   and the interface with the unoxidized bulk was well coher  served in the subscales, both in its original particle scattered shape   ent, as displayed in Fig. 20 c. Fibers were absent in the oxide layer   and at the fiber/matrix interface in the form of lamellae, Fig. 17 e.   and remained instead preserved in the first boride grains. No SiC  Then, the typical spinodal microstructure with C thin flakes was   depleted ZrB 2   layer was detected, diversely from what is generally   found moving out of the central valley both in ZrB 2   grains and in   reported for ZrB 2  -SiC bulk [52] suggesting that carbon might sup  ZrO 2   grains towards the outer surface, Fig. 17 b,d&e. Also, the EDS   press SiC active oxidation, possibly by preventing the formation of   elemental mapping recorded in the center of the valley confirmed   a stable SiO 2   outermost layer that varies the oxygen partial pres  only limited oxygen penetration and preservation of the boride   sure across the section.   matrix, fiber and SiC phase, Fig. 18 .   Moving to the short-fiber containing UHTC, it can be appreci  A notably different appearance was that of the FJ-LF-2 sam  ated that the FJ-SF model achieved 2930 K and its shape resulted   ple which achieved about 2400 K peak temperature and ended   accordingly the most compromised of the three tests. Given the   the test without meaningful shape and size variation. The sur  slight misalignment and the power of the jet, the surface looked   11   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 17. SEM images of the polished cross section of the FJ-LF-1 sample after test showing a) the overall section profile with magnified pictures as labeled and EDS spectra   of the mentioned phases.   Fig. 18. SEM image and EDS elemental mapping of the FJ-LF-1 sample taken from the “c” zone in Fig. 15 .   12   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 19. SEM images of the surface of the FJ-LF-2 sample after test showing a) the overall aspect with magnified views of the b) central and c) edge zones.   Fig. 20. SEM images of the polished cross section of the FJ-LF-2 sample after test showing a) the general overview with EDS elemental mapping, b) an enlarged view of the   outer oxide and c) of the oxide/boride interface.   13   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 21. a) Surface of the FJ-SF sample with enlarged views of different areas as indicated.   quite inhomogeneous depending on the position, Fig. 21 a. In this   3.2. Chamber insert tests   sample, all the eutectic compositions mentioned for the FJ-LF-1   sample might have formed, plus, the peak temperature was close   For the second set up of the experimental characterization cam  to that of ZrO 2   melting, which explains the effect of shear forces   paign on UHTCMC chamber inserts, two flat disks with the matrix   on the lateral material ablation. In the areas less hit by the jet,   based on ZrB 2   as major component and SiC as a minority phase   silica glass was found as well as ZrO 2   with generally coarse mor  have been manufactured and tested, one with long carbon fibers   phology, Fig. 21 b. In the center, round pits were found. The fact   and the other with chopped fibers uniformly dispersed into the   that ZrO 2   had a finer structure within these pits as compared to   matrix, Table 2 . Moreover, a classical C/SiC flat disk sample, ob  the outer region, Fig. 21 c, suggests that a cap of dense oxide was   tained by polymer infiltration and pyrolysis (PIP) technology, was   detached leaving exposed the underneath layer after violent gas re  tested as reference material. All the samples have been tested in   lease. The oxidation process of UHTC materials is actually featured   Test Condition “CI”: the measured oxidizer mass flow rate was   by accumulation of gases in the subscales and release in form of   40.4 g/s in all the tests, whereas the time-averaged combustion   scattered events with abrupt and vigorous explosion-like phenom  chamber pressure was about 8.6 bar.   ena, especially under atmospheres lacking of oxygen, which could   Starting from the C/SiC chamber insert, Fig. 23 shows the pic  instead promote partial healing in the first ablation stage [31] .   tures of the sample before and after test, where the surface ex  Then, moving to the most eroded region, the material removal was   posed to the flame is clearly observable, such as the enlargement   so important to reach the material core and bring to light the car  of the transversal section, whose diameter increased from the ini  bon fibers that remained unoxidized immersed in a thin ZrO 2  layer,   tial value of 15 mm to a final value of around 20.6 mm after the   Fig. 21 d.   test, Fig. 23 c. Moreover, also for what concerns the structural resis  Analysis of the cross section of the FJ-SF sample evidenced that   tance, although no cracks were detected, a delamination of parts of   about a half of the specimen height was consumed by the jet,  μm underwent oxidation in the   Fig. 22 a. However, less than 60   the first layer was found, as it can be seen from Fig. 23 b. The low   value of interlaminar shear strength of these CMCs is well known   central zone, as demonstrated by the EDS elemental mapping in   in the literature [53] .   Fig. 22 b. The oxide structure was composed of a coarse/melted   For what concerns the UHTCMC chamber inserts, Fig. 24 and   ZrO 2   scale which topped a fine grained one, both interrupted by   Fig. 25 show respectively pictures of the long-fiber and short  bubbles and silica pockets, Fig. 22 d. The interface oxide/boride was   fiber chamber inserts before and after test. After the test the CI-LF   in-built without relevant delamination phenomena. Likewise sam  chamber insert presented visible oxidation on the surface exposed   ple FJ-LF-2, right anchored to the oxide, the boride matrix was   to the reacting flow, whilst the internal hole surface did show nei  filled with SiC needles and platelets and the fibers underwent a   ther significant material erosion, nor structural failures, but only   sort of “second ultra-high temperature sintering” which changed   little deposit and surface roughening, Fig. 24 b&c. On the other side,   their profile and interface with the matrix and resulted in irregu  CI-SF sample was subjected to fatal structural cracks, which, al  lar edges with a crown of SiC platelets aligned to ZrB 2   elongated   though left the throat diameter unchanged, led to leakage of the   grains, Fig. 22 c&e.   combusting gasses determining critical damages also to the engine,   14   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 22. a) Cross section of the FJ-SF sample with magnification of the areas as indicated. b) Shows the SEM image with EDS elemental mapping in the central zone and   EDS spectra of the mentioned phases are reported below.   Fig. 23. Pictures of a chamber insert made of conventional C/SiC: a) before test; b), c) after test showing oxide spall-off and notable erosion of the central throat. The dotted   line in b) and c) marks the original throat diameter.   15   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   Fig. 24. Pictures of the CI-LF chamber insert a) before and b), c) after test showing no erosion of the central throat. Microstructural details of the pristine material are shown   in a).   Fig. 25. Pictures of the CI-SF chamber insert a) before and b), c) after test. Microstructural details of the pristine material are shown in a).   Fig. 25 . Anyway, it is worth to note that, besides the zones around   Samples based on a ZrB 2   - 10 vol% SiC matrix with continuous   the crack, the internal surfaces of the hole did not show significant   or chopped carbon fibers were sintered by either Hot Pressing or   material erosion.   Spark Plasma Sintering. Two long-fiber and one short-fiber samples   The diverse behavior of the chamber inserts can be ascribed   were tested in the free-jet configuration. One of the long fibers   to the thermal shock resistance of the starting material, i.e. in   samples displayed an outstanding erosion resistance, despite sur  the case long fiber, only when about 15% porosity was left after   face oxidation, while the other two, which reached surface temper  sintering no structural damage occurred. Accordingly, the mate  atures over 2800 K after a sudden temperature rise, were subjected   rial with short fibers (which provide weaker mechanical reinforce  to a more significant mass loss (erosion rates around 0.15 mm/s).   ment), having a fully dense matrix and whose fibers integrity had   However, all the UHTCMC materials demonstrated an excellent per  been compromised during sintering, behaved as a brittle ceramic   formance when compared to a reference graphite sample tested in   and shattered.   similar conditions.   Therefore, the results described above confirm the fact that the   Two chamber inserts (one with long and one with short car  ZrB 2  /SiC-based UHTCMC are characterized by outstanding erosion   bon fibers) were manufactured and tested. The long-fiber sam  resistance, provided that sufficient porosity is left to guarantee   ple survived the test without erosion nor structural damage, while   enough thermal shock tolerance.   4. Conclusions   This paper presented a range of experimental activities carried   out for the characterization of a new class of UHTCMCs for the de  velopment of near-zero erosion rocket engine components. To per  form a preliminary material screening and select the most suitable   candidates for manufacturing large and complex engine prototypes,   two dedicated test set-ups were developed, upgrading an existing   experimental bench for testing hybrid rocket engines. In the first   set-up (“Free-jet test”), small button-like samples were exposed to   the supersonic exhaust jet of a 200N-class hybrid rocket motor in   order to assess the capability of the material to withstand harsh   aero-thermo-chemical environment in moderate aerodynamic con  ditions, in a stagnation point configuration: heat flux in the order   of 15 MW/m  2 , temperature above 2500 K, pressure in the order of   3-4 bar, oxidizing and chemically reacting atmosphere. The second   set-up (“Chamber insert test”) was meant to assess the capability   of the test article (in the shape and size of an annular combustion   the short-fiber prototype was subjected to critical mechanical fail  ure. The different behavior was attributed to the sintering condi  tions, which resulted in a higher porosity of the long-fiber ma  terial, favouring an improved thermal shock resistance. Both test   articles performed better than reference C/SiC in terms of erosion   resistance.   Overall, the present test campaign demonstrated a promis  ing behavior of carbon-fiber reinforced ZrB 2  -based UHTCMCs in   a harsh environment representative of hybrid rocket applications,   with an excellent erosion resistance with respect to benchmark   state-of-art materials. The results suggest moreover that the man  ufacturing process needs to be carefully controlled in order to pro  vide the final components with optimal mechanical properties.   Declaration of Competing Interest   The authors declare that they have no known competing finan  cial interests or personal relationships that could have appeared to   influence the work reported in this paper.   chamber element) to withstand high thermo-mechanical stress at   CRediT authorship contribution statement   high pressures (order of 10 bar) in relevant aero-thermo-chemical   combustion environment (heat flux around 5 MW/m  2 ). The experi  Stefano Mungiguerra: Conceptualization, Methodology, Vali  mental observations included erosion rate estimation, surface tem  dation, Formal analysis, Investigation, Data curation, Writing   perature measurements by means of infrared equipment (pyrome  original draft, Visualization. Giuseppe D. Di Martino: Conceptual  ters, thermo-camera) and post-test microstructure analyses.   ization, Methodology, Validation, Investigation, Writing original   16   \\x0c', 'S. Mungiguerra, G.D. Di Martino, R. Savino et al.   International Journal of Heat and Mass Transfer 163 (2020) 120492   draft, Visualization. Raffaele Savino: Conceptualization, Method  ology, Supervision, Funding acquisition. Luca Zoli: Methodology,   Investigation, Resources, Writing review & editing. Laura Silve  stroni: Methodology, Investigation, Resources, Writing review &   editing. Diletta Sciti: Supervision, Funding acquisition, Project ad  ministration, Writing review & editing.   Acknowledgements   This study has received funding by the European Union’s Hori  zon2020 research and innovation program under the research   project C  3 HARME with Grant Agreement No. 685594.   The authors wish to thank Prof. Claudio Leone and Dr. Silvio   Genna (CIRTIBS Research Center) for technical support in realizing   the microscopic pictures of the samples.   The authors would like to acknowledge A. Schoberth from AIR  BUS CRT for providing the reference C/SiC material for chamber in  sert testing; M. Lagos from TECNALIA for spark plasma sintering of   samples FJ-LF-1 and FJ-SF; and S. Rivera from Nanoker Research for   spark plasma sintering of CI-SF chamber insert.   Supplementary materials   Supplementary material associated with this article can be   found, in the online version, at doi:10.1016/j.ijheatmasstransfer.   2020.120492 .   References   [1] R. Hickman , T. Mc Kechnie , A. Agarwal , Net shape fabrication of high tempera  ture materials for rocket engine components, 37th Joint Propulsion Conference   and Exhibit, 2001 AIAA paper 2001-3435 .   [2] Johnston, J.R., Signorelli, R.A., Freche, J.C., “Performances of rocket nozzle ma  terial with several solid propellants,” NASA technical note 3428.3, 1966.   [3] W. Chen , Numerical analyses of ablative behavior of C/C composite materials,   Int J Heat Mass Transf 95 (2016) 720-726 .   [4] P. Thakre , V. Yang , Chemical erosion of carbon-carbon/graphite nozzles in   solid-propellant rocket motors, J Propuls. 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},{
  "_id": 28,
  "PDF": "coatings-07-00111.pdf",
  "Text": "['Article  Thermal Analysis of Tantalum Carbide-Hafnium Carbide Solid Solutions from Room Temperature to 1400  C  Cheng Zhang, Archana Loganathan, Benjamin Boesl and Arvind Agarwal *  Plasma Forming Laboratory, Department of Mechanical and Materials Engineering, Florida International  University, Miami, 33139 FL, USA; czhan009@fiu.edu (C.Z.); aloga006@fiu.edu (A.L.); bboesl@fiu.edu (B.B.) * Correspondence: agarwala@fiu.edu; Tel.: +1-305-348-1701  Received: 5 June 2017; Accepted: 25 July 2017; Published: 28 July 2017  Abstract: The thermogravimetric analysis on TaC, HfC, and their solid solutions has been carried out in air up to 1400  C. Three solid solution compositions have been chosen: 80TaC-20 vol % HfC (T80H20), 50TaC-50 vol % HfC (T50H50), and 20TaC-80 vol % HfC (T20H80), in addition to pure  TaC and HfC. Solid solutions exhibit better oxidation resistance than the pure carbides. The onset of oxidation is delayed in solid solutions from 750  C for pure TaC, to 940  C for the T50H50 sample. Moreover, T50H50 samples display the highest resistance to oxidation with the retention of the initial  carbides. The oxide scale formed on the T50H50 sample displays mechanical integrity to prevent  the oxidation of the underlying carbide solid solution. The improved oxidation resistance of the solid solution is attributed to the reaction between Ta2O5 and HfC, which stabilizes the volume changes induced by the formation of Ta2O5 and diminishes the generation of gaseous products. Also, the formation of solid solutions disturbs the atomic arrangement inside the lattice, which delays the  reaction between Ta and O. Both of these mechanisms lead to the improved oxidation resistances of  TaC-HfC solid solutions.  Keywords: tantalum carbide; hafnium carbide; solid solutions; oxidation; thermogravimetric analysis  1. Introduction  The interest in tantalum carbide (TaC) and hafnium carbide (HfC) has been growing in recent years  due to their extremely high melting points, high hardness, and elastic moduli, and more importantly,  their ability to form solid solutions [1-4]. The major applications of these two carbides are leading  edges of reentry vehicles and lining materials for rocket thrusters. In both cases, excellent oxidation resistance is required. However, gaseous products like CO and CO2 are inevitably formed during oxidation, which leads to porous oxide scales that delaminate and spall. The major oxide of TaC is Ta2O5 , which has a melting point of ~1900  C, lower than the desired application temperature of 2000  C or more [5-7]. As a result, the resultant oxide would melt and lose its structural integrity, and fail catastrophically. To reduce the gaseous products as well as retain the integrity of oxide scales under extremely high temperatures, Hafnium diboride (HfB2 ) and its composites have been investigated as promising candidate materials for use on next-generation hypersonic vehicles [8,9]. During oxidation, HfB2 forms a solid scaffold-like structure that mainly consists of HfO2 and molten B2O3 infiltrated between the HfO2 . The resultant oxide scale is dense and crack-free, which provides exceptional the B2O3 starts to evaporate around 700  C and therefore its oxidation resistance. Unfortunately, protection of the underlying materials is lost. The SiC addition was used to stabilize the B2O3 by forming a borosilicate glass phase, which increases the onset evaporation temperature to 1400  C. However, 1400  C is still not high enough to withstand higher application temperatures of 2000  C or more.  Coatings 2017, 7, 111; doi:10.3390/coatings7080111  www.mdpi.com/journal/coatings  coatings\\x0c', 'Coatings 2017, 7, 111  2 of 9  The studies on the solid solutions of TaC-HfC began with the discovery of a TaC0.8HfC0.2 phase that possesses the highest melting point (~4000  C) of known substances [10]. Preliminary oxidation studies have been carried out on TaC0.8HfC0.2 and HfC-rich compositions, but no improvement in the oxidation behavior was observed compared to pure carbides [11-14]. Additionally, sintering aids were  inevitable in those studies, which introduced secondary phases that clouded the understanding of oxidation behaviors. Although TaC and HfC can form solid solutions above 887  C in all compositions, as shown in the phase diagram in Figure 1, oxidation studies on TaC-HfC solid solutions have barely  been investigated.  Figure 1. Phase diagram of TaC and HfC [15]. Copyright 2013 Elsevier.  Recently, Cedillos-Barraza et al. as well as our research group sintered TaC-HfC solid solutions  without sintering additions by spark plasma sintering (SPS) [16,17]. The compositions cover the full spectrum of TaC-HfC solid solutions, and both studies noticed that TaC0.5HfC0.5 has the highest hardness and elastic modulus among the solid solutions. Our group conducted oxidation testing using a plasma jet by exposing these solid solutions and pure carbides to a temperature of ~3000  C at a gas flow rate of sonic speed [18]. In general, the solid solutions showed better oxidation resistance than the pure carbides. The best oxidation resistance was found in the TaC0.5HfC0.5 composition. After 300 s of exposure to such extreme oxidation conditions, the thickness of the oxide scale in TaC0.5HfC0.5 was only 28 µm, which is 1/6 and 1/10 of the oxide scale thickness in pure HfC and TaC, respectively [18]. The improved oxidation mechanism was explained by a newly formed Hf6Ta2O17 phase. More importantly, we found a similar dense solid scaffold and liquid phase structure as reported in HfB2 -SiC/HfB2 -B4C systems that protect the underlying materials [18]. In the case of TaC-HfC solid solutions, the solid scaffold consists of HfO2 and Hf6Ta2O17 , and the liquid phase is made of molten Ta2O5 . Compared to the B2O3 and borosilicate phase in the diboride system, molten Ta2O5 is a much more stable phase with a higher melting point of 1900  C. Hence, the carbide solid solutions exhibit exceptional oxidation resistance.  One question arises after the investigation on the plasma jet oxidation behavior of the carbide solid solutions: How would the carbide solid solutions behave below 1800  C, where the temperature is not high enough to melt the resultant Ta2O5 ? To address this question, we sought to understand the carbide solid solutions from room temperature to 1400  C using the oxidation behavior of thermogravimetric analysis (TGA). Five samples, namely pure TaC, 80TaC-20 vol % HfC (T80H20),  50TaC-50 vol % HfC (T50H50), 20TaC-80 vol % HfC (T20H80), and pure HfC, were chosen. A detailed  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa02\\xa0of\\xa09\\xa0\\xa0The\\xa0studies\\xa0on\\xa0the\\xa0solid\\xa0solutions\\xa0of\\xa0TaC‐HfC\\xa0began\\xa0with\\xa0the\\xa0discovery\\xa0of\\xa0a\\xa0TaC0.8HfC0.2\\xa0phase\\xa0that\\xa0possesses\\xa0the\\xa0highest\\xa0melting\\xa0point\\xa0(~4000\\xa0°C)\\xa0of\\xa0known\\xa0substances\\xa0[10].\\xa0Preliminary\\xa0oxidation\\xa0studies\\xa0have\\xa0been\\xa0carried\\xa0out\\xa0on\\xa0TaC0.8HfC0.2\\xa0and\\xa0HfC‐rich\\xa0compositions,\\xa0but\\xa0no\\xa0improvement\\xa0in\\xa0the\\xa0oxidation\\xa0behavior\\xa0was\\xa0observed\\xa0compared\\xa0to\\xa0pure\\xa0carbides\\xa0[11-14].\\xa0Additionally,\\xa0sintering\\xa0aids\\xa0were\\xa0inevitable\\xa0in\\xa0those\\xa0studies,\\xa0which\\xa0introduced\\xa0secondary\\xa0phases\\xa0that\\xa0clouded\\xa0the\\xa0understanding\\xa0of\\xa0oxidation\\xa0behaviors.\\xa0Although\\xa0TaC\\xa0and\\xa0HfC\\xa0can\\xa0form\\xa0solid\\xa0solutions\\xa0above\\xa0887\\xa0°C\\xa0in\\xa0all\\xa0compositions,\\xa0as\\xa0shown\\xa0in\\xa0the\\xa0phase\\xa0diagram\\xa0in\\xa0Figure\\xa01,\\xa0oxidation\\xa0studies\\xa0on\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0have\\xa0barely\\xa0been\\xa0investigated.\\xa0\\xa0Figure\\xa01.\\xa0Phase\\xa0diagram\\xa0of\\xa0TaC\\xa0and\\xa0HfC\\xa0[15].\\xa0Copyright\\xa02013\\xa0Elsevier.\\xa0Recently,\\xa0Cedillos‐Barraza\\xa0et\\xa0al.\\xa0as\\xa0well\\xa0as\\xa0our\\xa0research\\xa0group\\xa0sintered\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0without\\xa0sintering\\xa0additions\\xa0by\\xa0spark\\xa0plasma\\xa0sintering\\xa0(SPS)\\xa0[16,17].\\xa0The\\xa0compositions\\xa0cover\\xa0the\\xa0full\\xa0spectrum\\xa0of\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0and\\xa0both\\xa0studies\\xa0noticed\\xa0that\\xa0TaC0.5HfC0.5\\xa0has\\xa0the\\xa0highest\\xa0hardness\\xa0and\\xa0elastic\\xa0modulus\\xa0among\\xa0the\\xa0solid\\xa0solutions.\\xa0Our\\xa0group\\xa0conducted\\xa0oxidation\\xa0testing\\xa0using\\xa0a\\xa0plasma\\xa0jet\\xa0by\\xa0exposing\\xa0these\\xa0solid\\xa0solutions\\xa0and\\xa0pure\\xa0carbides\\xa0to\\xa0a\\xa0temperature\\xa0of\\xa0~3000\\xa0°C\\xa0at\\xa0a\\xa0gas\\xa0flow\\xa0rate\\xa0of\\xa0sonic\\xa0speed\\xa0[18].\\xa0In\\xa0general,\\xa0the\\xa0solid\\xa0solutions\\xa0showed\\xa0better\\xa0oxidation\\xa0resistance\\xa0than\\xa0the\\xa0pure\\xa0carbides.\\xa0The\\xa0best\\xa0oxidation\\xa0resistance\\xa0was\\xa0found\\xa0in\\xa0the\\xa0TaC0.5HfC0.5\\xa0composition.\\xa0\\xa0After\\xa0300\\xa0s\\xa0of\\xa0exposure\\xa0to\\xa0such\\xa0extreme\\xa0oxidation\\xa0conditions,\\xa0the\\xa0thickness\\xa0of\\xa0the\\xa0oxide\\xa0scale\\xa0in\\xa0TaC0.5HfC0.5\\xa0was\\xa0only\\xa028\\xa0μm,\\xa0which\\xa0is\\xa01/6\\xa0and\\xa01/10\\xa0of\\xa0the\\xa0oxide\\xa0scale\\xa0thickness\\xa0in\\xa0pure\\xa0HfC\\xa0and\\xa0TaC,\\xa0respectively\\xa0[18].\\xa0The\\xa0improved\\xa0oxidation\\xa0mechanism\\xa0was\\xa0explained\\xa0by\\xa0a\\xa0newly\\xa0formed\\xa0Hf6Ta2O17\\xa0phase.\\xa0More\\xa0importantly,\\xa0we\\xa0found\\xa0a\\xa0similar\\xa0dense\\xa0solid\\xa0scaffold\\xa0and\\xa0liquid\\xa0phase\\xa0structure\\xa0as\\xa0reported\\xa0in\\xa0HfB2‐SiC/HfB2‐B4C\\xa0systems\\xa0that\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0[18].\\xa0In\\xa0the\\xa0case\\xa0of\\xa0\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0the\\xa0solid\\xa0scaffold\\xa0consists\\xa0of\\xa0HfO2\\xa0and\\xa0Hf6Ta2O17,\\xa0and\\xa0the\\xa0liquid\\xa0phase\\xa0is\\xa0made\\xa0of\\xa0molten\\xa0Ta2O5.\\xa0Compared\\xa0to\\xa0the\\xa0B2O3\\xa0and\\xa0borosilicate\\xa0phase\\xa0in\\xa0the\\xa0diboride\\xa0system,\\xa0molten\\xa0Ta2O5\\xa0is\\xa0a\\xa0much\\xa0more\\xa0stable\\xa0phase\\xa0with\\xa0a\\xa0higher\\xa0melting\\xa0point\\xa0of\\xa01900\\xa0°C.\\xa0Hence,\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0exhibit\\xa0exceptional\\xa0oxidation\\xa0resistance.\\xa0One\\xa0question\\xa0arises\\xa0after\\xa0the\\xa0investigation\\xa0on\\xa0the\\xa0plasma\\xa0jet\\xa0oxidation\\xa0behavior\\xa0of\\xa0the\\xa0carbide\\xa0solid\\xa0solutions:\\xa0How\\xa0would\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0behave\\xa0below\\xa01800\\xa0°C,\\xa0where\\xa0the\\xa0temperature\\xa0is\\xa0not\\xa0high\\xa0enough\\xa0to\\xa0melt\\xa0the\\xa0resultant\\xa0Ta2O5?\\xa0To\\xa0address\\xa0this\\xa0question,\\xa0we\\xa0sought\\xa0to\\xa0understand\\xa0the\\xa0oxidation\\xa0behavior\\xa0of\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0from\\xa0room\\xa0temperature\\xa0to\\xa01400\\xa0°C\\xa0using\\xa0thermogravimetric\\xa0analysis\\xa0(TGA).\\xa0Five\\xa0samples,\\xa0namely\\xa0pure\\xa0TaC,\\xa080TaC‐20\\xa0vol\\xa0%\\xa0HfC\\xa0(T80H20),\\xa050TaC‐50\\xa0vol\\xa0%\\xa0HfC\\xa0(T50H50),\\xa020TaC‐80\\xa0vol\\xa0%\\xa0HfC\\xa0(T20H80),\\xa0and\\xa0pure\\xa0HfC,\\xa0were\\xa0chosen.\\xa0A\\xa0detailed\\xa0\\x0c', 'Coatings 2017, 7, 111  3 of 9  analysis of the oxidation behavior is carried out in the present study using TGA followed by scanning  electron microscopy (SEM).  2. Experimental Details  2.1. Materials  Commercial TaC powder (Inframat Advanced Materials LLC, Manchester, CT, USA) and hafnium  carbide powder (Materion LLC, Cleveland, OH, USA) were used as starting powders. Powders for  solid solutions were mixed by a high-energy vibratory ball milling machine (Across International LLC,  Livingston, NJ, USA) according to their stoichiometric ratio. Pure powders were milled for one hour  separately in a tungsten carbide (WC) jar to breakdown the agglomeration. Subsequently, TaC and  HfC powders were mixed for another hour. The ball to powder ratio was 1:3, using a 6-mm diameter  WC ball. The mixed TaC-HfC powders were consolidated by a spark plasma sintering (SPS) machine (Model 10-4, Thermal Technologies, Santa Rose, CA, USA). The powders were sintered at 1850  C with a heating rate of 100  C/min and a maximum uniaxial pressure of 60 MPa. The holding time was 10 min to ensure the densification. The environment in the vacuum was set at a pressure of 4 Pa.  The details of the processing can be found in our previous study [17].  2.2. TGA Testing and Post-Oxidation Characterization  The TGA oxidation testing was conducted on a small portion (~30 mg) of the sintered pellets.  A thermogravimetric analysis (TGA) analyzer (SDT-Q600, TA Instruments, New Castle, DE, USA)  was used to evaluate the oxidation performance of TaC, HfC, and TaC-HfC solid solution samples. Samples were tested in the air at a heating rate of 5  C/min. The maximum temperature was 1400  C for all samples. The morphologies of the post-oxidation samples were examined by a field emission  SEM (JSM-6330F, JEOL Ltd., Tokyo, Japan).  3. Results and Discussions  3.1. Microstructure and Phases in Sintered TaC, HfC, and TaC-HfC Solid Solutions  The detailed characterization results of  the microstructures  and phases  formed in spark  plasma-sintered TaC-HfC solid solutions were published in our previous paper [17]. For the reader ’s  convenience and the sake of completeness, a summary of the key results is listed in Table 1. All five  samples had high densities varying between 97% and 99%. With the addition of HfC, the samples’  densification increases and the highest densification is found in T20H80 sample. The average grain size  also decreased with the HfC additions. The lattice parameters of the formed solid solutions matched  the theoretical values calculated according to Vegard’s Law. [17]  Table 1. Basic characterizations of the TaC, HfC, and TaC-HfC solid solutions.  Name  Pure TaC T80H20 T50H50 T20H80 Pure HfC  Pellet Density (×103 kg/m3 )  Densification (%)  14.14 13.85 13.26 12.68 12.21  96.7 97.8 98.2 98.8 98.5  Average Grain Size (µm) 6.8 ± 1.4 6.2 ± 2.1 3.8 ± 1.2 3.1 ± 1.1 2.3 ± 0.7  3.2. Macro State Morphology of Post-Oxidation TaC-HfC Solid Solutions  The overall qualitative oxidation resistance of TaC, HfC, and their solid solutions can be inferred  by the morphology of the post-oxidation samples. Appearances of the post-oxidation samples from  the TGA testing are shown in Figure 2. The pure TaC sample showed the worst oxidation resistance,  as it turned into a powdery form with no mechanical integrity (Figure 2a). The oxidized pure HfC,  \\x0c', 'Coatings 2017, 7, 111  4 of 9  on the other hand, displayed a much better oxidation resistance. Structural integrity can still be seen  even though the oxidized sample broke into several pieces (Figure 2b). The post-oxidation samples’  appearances of T80H20 and T20H80 is the combination of oxidation morphology exhibited by pure  TaC and HfC. In the oxidized T80H20 sample (Figure 2c), a large amount of powder is noticed with  a few solid broken pieces. The appearance of the post-oxidation T20H80 (Figure 2d) is analogous to  pure HfC with a small amount of powder. T50H50 is the only sample which does not show significant  spallation and delamination. This suggests that outer layer oxide scale has good mechanical integrity  and can protect the underlying carbide solid solution.  Figure 2. Post-oxidation samples: (a) Pure TaC; (b) Pure HfC; (c) T80H20; (d) T20H80; and (e) T50H50.  3.3. Mass Change during Thermogravimetric Analysis of Carbide Solid Solutions  The weight change curves of TaC-HfC solid solutions are presented in Figure 3 as the degree  of oxidation (α) with the onset oxidation temperatures for five samples. The degree of oxidation (α)  is defined as the ratio of the measured weight change over the theoretical weight change at 100%  conversion. The degree of oxidation (α) is calculated by the following equation:  ∝=  ∆m  ∆m∞  =  mins − mi  mt − mi  where ∆m is the measured weight change, ∆m∞ is the theoretical weight change, mins measured weight, mi is the initial weight, and mt is the theoretical weight. The theoretical weight change is based on the below reactions:  is the real-time  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa04\\xa0of\\xa09\\xa0\\xa0even\\xa0though\\xa0the\\xa0oxidized\\xa0sample\\xa0broke\\xa0into\\xa0several\\xa0pieces\\xa0(Figure\\xa02b).\\xa0The\\xa0post‐oxidation\\xa0samples’\\xa0appearances\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0is\\xa0the\\xa0combination\\xa0of\\xa0oxidation\\xa0morphology\\xa0exhibited\\xa0by\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0In\\xa0the\\xa0oxidized\\xa0T80H20\\xa0sample\\xa0(Figure\\xa02c),\\xa0a\\xa0large\\xa0amount\\xa0of\\xa0powder\\xa0is\\xa0noticed\\xa0with\\xa0a\\xa0few\\xa0solid\\xa0broken\\xa0pieces.\\xa0The\\xa0appearance\\xa0of\\xa0the\\xa0post‐oxidation\\xa0T20H80\\xa0(Figure\\xa02d)\\xa0is\\xa0analogous\\xa0to\\xa0pure\\xa0HfC\\xa0with\\xa0a\\xa0small\\xa0amount\\xa0of\\xa0powder.\\xa0T50H50\\xa0is\\xa0the\\xa0only\\xa0sample\\xa0which\\xa0does\\xa0not\\xa0show\\xa0significant\\xa0spallation\\xa0and\\xa0delamination.\\xa0This\\xa0suggests\\xa0that\\xa0outer\\xa0layer\\xa0oxide\\xa0scale\\xa0has\\xa0good\\xa0mechanical\\xa0integrity\\xa0and\\xa0can\\xa0protect\\xa0the\\xa0underlying\\xa0carbide\\xa0solid\\xa0solution.\\xa0\\xa0Figure\\xa02.\\xa0Post‐oxidation\\xa0samples:\\xa0(a)\\xa0Pure\\xa0TaC;\\xa0(b)\\xa0Pure\\xa0HfC;\\xa0(c)\\xa0T80H20;\\xa0(d)\\xa0T20H80;\\xa0and\\xa0(e)\\xa0T50H50.\\xa03.3.\\xa0Mass\\xa0Change\\xa0during\\xa0Thermogravimetric\\xa0Analysis\\xa0of\\xa0Carbide\\xa0Solid\\xa0Solutions\\xa0The\\xa0weight\\xa0change\\xa0curves\\xa0of\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0are\\xa0presented\\xa0in\\xa0Figure\\xa03\\xa0as\\xa0the\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0with\\xa0the\\xa0onset\\xa0oxidation\\xa0temperatures\\xa0for\\xa0five\\xa0samples.\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0is\\xa0defined\\xa0as\\xa0the\\xa0ratio\\xa0of\\xa0the\\xa0measured\\xa0weight\\xa0change\\xa0over\\xa0the\\xa0theoretical\\xa0weight\\xa0change\\xa0at\\xa0100%\\xa0conversion.\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0is\\xa0calculated\\xa0by\\xa0the\\xa0following\\xa0equation:\\xa0∝(cid:3404)∆(cid:1865)∆(cid:1865)(cid:2998)(cid:3404)(cid:1865)(cid:3036)(cid:3041)(cid:3046)(cid:3398)(cid:1865)(cid:3036)(cid:1865)(cid:3047)(cid:3398)(cid:1865)(cid:3036)\\xa0\\xa0where\\xa0∆(cid:1865)\\xa0is\\xa0the\\xa0measured\\xa0weight\\xa0change,\\xa0∆(cid:1865)(cid:2998)\\xa0is\\xa0the\\xa0theoretical\\xa0weight\\xa0change,\\xa0(cid:1865)(cid:3036)(cid:3041)(cid:3046)\\xa0is\\xa0the\\xa0\\xa0real‐time\\xa0measured\\xa0weight,\\xa0(cid:1865)(cid:3036)\\xa0is\\xa0the\\xa0initial\\xa0weight,\\xa0and\\xa0(cid:1865)(cid:3047)\\xa0is\\xa0the\\xa0theoretical\\xa0weight.\\xa0The\\xa0theoretical\\xa0weight\\xa0change\\xa0is\\xa0based\\xa0on\\xa0the\\xa0below\\xa0reactions:\\xa02TaC(cid:3397)92(cid:3415)O(cid:2870)→Ta(cid:2870)O(cid:2873)(cid:3397)2CO(cid:2870)\\xa0(1)\\xa0\\x0c', 'Coatings 2017, 7, 111  9  O2 → Ta2O5 + 2CO2  2TaC +  2 HfC + O2 → HfO2 + CO2  5 of 9  (1)  (2)  Figure 3. TG curve of TaC, HfC, and TaC-HfC solid solutions.  The onset oxidation temperature is defined in Figure 3 when the degree of oxidation increases sharply. Pure TaC starts to oxidize around 750  C, followed by pure HfC which starts around 800  C. The solid solution samples showed improved oxidation resistance, as the oxidation processes have  been delayed compared to the pure carbides, as shown in Figure 2. T80H20 began its oxidation process at 850  C, and T20H80 started around 900  C. Further delay was observed in T50H50, where the onset oxidation temperature was around 940  C. The oxidation process of pure TaC matches the literature description [19,20]. The degree of oxidation increases sharply around 750  C and is completed near 950  C. The transformation from TaC to Ta2O5 involves tremendous volume changes. According to the Pilling-Bedworth (PB) ratio theory, the PB ratio for Ta to Ta2O5 is 2.5 [21]. Such large volume mismatch leads to spallation and delamination. Without any anchoring structure, the resultant oxide will lose its mechanical integrity,  and thus it cannot protect the underlying materials. Consequently, the final product mainly consists of fine powder, as shown in Figure 2a. One spike at 780  C is observed in the TaC oxidation curve in Figure 3. The measured temperature dropped during the oxidation process. The morphology of  the oxidation product suggests that the oxide would delaminate and expose unreacted TaC during  oxidation. The unreacted TaC has a lower temperature than the surface of the delaminated oxide,  so the measured temperature dropped. Figure 3 shows that oxidation process of HfC starts at 800  C. Comparing to the oxidation of pure TaC, the weight increases gradually instead of exhibiting a sharp increase, as evident from the slope.  Such behavior is in accordance with the literature description of the oxidation behavior of HfC [12,22].  HfC has the ability to absorb oxygen without transforming into oxide. The dissolved oxygen sits on the carbon vacancies and forms an oxy-carbide layer. The formed HfO2 , unlike Ta2O5 , does not experience much volume change from the HfC. The PB ratio is only 1.7 [21], which explains the mechanically stable morphology of HfO2 in Figure 2b. The morphologies of the oxidized TaC and HfC are studied by SEM and shown in Figure 4.  The grain size of oxidized TaC is larger than oxidized HfC. Distinct grains can be spotted in the  oxidized HfC sample, but the surface of oxidized TaC is relatively smooth with coalesced grains due  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa05\\xa0of\\xa09\\xa0\\xa0HfC(cid:3397)O(cid:2870)→HfO(cid:2870)(cid:3397)CO(cid:2870)\\xa0(2)\\xa0The\\xa0onset\\xa0oxidation\\xa0temperature\\xa0is\\xa0defined\\xa0in\\xa0Figure\\xa03\\xa0when\\xa0the\\xa0degree\\xa0of\\xa0oxidation\\xa0increases\\xa0sharply.\\xa0Pure\\xa0TaC\\xa0starts\\xa0to\\xa0oxidize\\xa0around\\xa0750\\xa0°C,\\xa0followed\\xa0by\\xa0pure\\xa0HfC\\xa0which\\xa0starts\\xa0around\\xa0800\\xa0°C.\\xa0The\\xa0solid\\xa0solution\\xa0samples\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance,\\xa0as\\xa0the\\xa0oxidation\\xa0processes\\xa0have\\xa0been\\xa0delayed\\xa0compared\\xa0to\\xa0the\\xa0pure\\xa0carbides,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa02.\\xa0T80H20\\xa0began\\xa0its\\xa0oxidation\\xa0process\\xa0at\\xa0850\\xa0°C,\\xa0and\\xa0T20H80\\xa0started\\xa0around\\xa0900\\xa0°C.\\xa0Further\\xa0delay\\xa0was\\xa0observed\\xa0in\\xa0T50H50,\\xa0where\\xa0the\\xa0onset\\xa0oxidation\\xa0temperature\\xa0was\\xa0around\\xa0940\\xa0°C.\\xa0\\xa0Figure\\xa03.\\xa0TG\\xa0curve\\xa0of\\xa0TaC,\\xa0HfC,\\xa0and\\xa0TaC‐HfC\\xa0solid\\xa0solutions.\\xa0The\\xa0oxidation\\xa0process\\xa0of\\xa0pure\\xa0TaC\\xa0matches\\xa0the\\xa0literature\\xa0description\\xa0[19,20].\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0increases\\xa0sharply\\xa0around\\xa0750\\xa0°C\\xa0and\\xa0is\\xa0completed\\xa0near\\xa0950\\xa0°C.\\xa0The\\xa0transformation\\xa0from\\xa0TaC\\xa0to\\xa0Ta2O5\\xa0involves\\xa0tremendous\\xa0volume\\xa0changes.\\xa0According\\xa0to\\xa0the\\xa0Pilling‐Bedworth\\xa0(PB)\\xa0ratio\\xa0theory,\\xa0the\\xa0PB\\xa0ratio\\xa0for\\xa0Ta\\xa0to\\xa0Ta2O5\\xa0is\\xa02.5\\xa0[21].\\xa0Such\\xa0large\\xa0volume\\xa0mismatch\\xa0leads\\xa0to\\xa0spallation\\xa0and\\xa0delamination.\\xa0Without\\xa0any\\xa0anchoring\\xa0structure,\\xa0the\\xa0resultant\\xa0oxide\\xa0will\\xa0lose\\xa0its\\xa0mechanical\\xa0integrity,\\xa0and\\xa0thus\\xa0it\\xa0cannot\\xa0protect\\xa0the\\xa0underlying\\xa0materials.\\xa0Consequently,\\xa0the\\xa0final\\xa0product\\xa0mainly\\xa0consists\\xa0of\\xa0fine\\xa0powder,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa02a.\\xa0One\\xa0spike\\xa0at\\xa0780\\xa0°C\\xa0is\\xa0observed\\xa0in\\xa0the\\xa0TaC\\xa0oxidation\\xa0curve\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0measured\\xa0temperature\\xa0dropped\\xa0during\\xa0the\\xa0oxidation\\xa0process.\\xa0The\\xa0morphology\\xa0of\\xa0the\\xa0oxidation\\xa0product\\xa0suggests\\xa0that\\xa0the\\xa0oxide\\xa0would\\xa0delaminate\\xa0and\\xa0expose\\xa0unreacted\\xa0TaC\\xa0during\\xa0oxidation.\\xa0The\\xa0unreacted\\xa0TaC\\xa0has\\xa0a\\xa0lower\\xa0temperature\\xa0than\\xa0the\\xa0surface\\xa0of\\xa0the\\xa0delaminated\\xa0oxide,\\xa0so\\xa0the\\xa0measured\\xa0temperature\\xa0dropped.\\xa0Figure\\xa03\\xa0shows\\xa0that\\xa0oxidation\\xa0process\\xa0of\\xa0HfC\\xa0starts\\xa0at\\xa0800\\xa0°C.\\xa0Comparing\\xa0to\\xa0the\\xa0oxidation\\xa0of\\xa0pure\\xa0TaC,\\xa0the\\xa0weight\\xa0increases\\xa0gradually\\xa0instead\\xa0of\\xa0exhibiting\\xa0a\\xa0sharp\\xa0increase,\\xa0as\\xa0evident\\xa0from\\xa0the\\xa0slope.\\xa0Such\\xa0behavior\\xa0is\\xa0in\\xa0accordance\\xa0with\\xa0the\\xa0literature\\xa0description\\xa0of\\xa0the\\xa0oxidation\\xa0behavior\\xa0of\\xa0HfC\\xa0[12,22].\\xa0HfC\\xa0has\\xa0the\\xa0ability\\xa0to\\xa0absorb\\xa0oxygen\\xa0without\\xa0transforming\\xa0into\\xa0oxide.\\xa0The\\xa0dissolved\\xa0oxygen\\xa0sits\\xa0on\\xa0the\\xa0carbon\\xa0vacancies\\xa0and\\xa0forms\\xa0an\\xa0oxy‐carbide\\xa0layer.\\xa0The\\xa0formed\\xa0HfO2,\\xa0unlike\\xa0Ta2O5,\\xa0does\\xa0not\\xa0experience\\xa0much\\xa0volume\\xa0change\\xa0from\\xa0the\\xa0HfC.\\xa0The\\xa0PB\\xa0ratio\\xa0is\\xa0only\\xa01.7\\xa0[21],\\xa0which\\xa0explains\\xa0the\\xa0mechanically\\xa0stable\\xa0morphology\\xa0of\\xa0HfO2\\xa0in\\xa0Figure\\xa02b.\\xa0The\\xa0morphologies\\xa0of\\xa0the\\xa0oxidized\\xa0TaC\\xa0and\\xa0HfC\\xa0are\\xa0studied\\xa0by\\xa0SEM\\xa0and\\xa0shown\\xa0in\\xa0Figure\\xa04.\\xa0The\\xa0grain\\xa0size\\xa0of\\xa0oxidized\\xa0TaC\\xa0is\\xa0larger\\xa0than\\xa0oxidized\\xa0HfC.\\xa0Distinct\\xa0grains\\xa0can\\xa0be\\xa0spotted\\xa0in\\xa0the\\xa0oxidized\\xa0HfC\\xa0sample,\\xa0but\\xa0the\\xa0surface\\xa0of\\xa0oxidized\\xa0TaC\\xa0is\\xa0relatively\\xa0smooth\\xa0with\\xa0coalesced\\xa0grains\\xa0due\\xa0to\\xa0localized\\xa0sintering.\\xa0Conventionally,\\xa0sintering\\xa0occurs\\xa0at\\xa00.6\\xa0Tm\\xa0(where\\xa0Tm\\xa0is\\xa0the\\xa0melting\\xa0point).\\xa0The\\xa0melting\\xa0point\\xa0(Tm)\\xa0of\\xa0Ta2O5\\xa0is\\xa01872\\xa0°C;\\xa0hence,\\xa0the\\xa0sintering\\xa0of\\xa0Ta2O5\\xa0will\\xa0start\\xa0around\\xa01250\\xa0°C.\\xa0\\xa0No\\xa0sintering\\xa0is\\xa0expected\\xa0in\\xa0HfO2,\\xa0as\\xa0its\\xa0melting\\xa0point\\xa0is\\xa0around\\xa02800\\xa0°C\\xa0[5].\\xa0The\\xa0pores\\xa0in\\xa0the\\xa0oxidized\\xa0HfC\\xa0are\\xa0from\\xa0the\\xa0formation\\xa0of\\xa0a\\xa0gaseous\\xa0product\\xa0during\\xa0oxidation.\\xa0\\x0c', 'Coatings 2017, 7, 111  6 of 9  to localized sintering. Conventionally, sintering occurs at 0.6 Tm (where Tm is the melting point). The melting point (Tm ) of Ta2O5 is 1872  C; hence, the sintering of Ta2O5 will start around 1250  C. No sintering is expected in HfO2 , as its melting point is around 2800  C [5]. The pores in the oxidized HfC are from the formation of a gaseous product during oxidation.  Figure 4. Post-oxidation SEM images of (a) pure TaC and (b) pure HfC.  Although the post-oxidation samples of T80H20 and T20H80 do not have sufficient mechanical  integrity to protect  the underlying materials  (Figure  2c,d),  these  two solid solutions  showed  improved oxidation resistance by delaying the onset of oxidation temperature, as shown in Figure 3. The oxidation onset temperatures of T80H20 and T20H80 are 850 and 900  C, respectively, which is significantly delayed as compared to pure TaC and HfC. However, both T80H20 and T20H80 still  reached near 100% oxidation. The T50H50 solid solution displayed the best oxidation resistance. Not only did the onset temperature for oxidation increase to near 940  C for the T50H50 samples, only 60% of the oxidation was completed when the temperature reached 1400  C. To further understand the mechanism of improved oxidation resistance in the solid solution  samples, especially in the T50H50 sample,  the morphologies of the post-oxidation samples of the  solid solutions are investigated by SEM and presented in Figure 5. As described earlier, the oxide  morphologies of pure TaC and HfC (Figure 4) are highly porous due to gaseous products resulting  in prominent volume change. The key factor in improving the oxidation resistance of the TaC-HfC solid solutions is to suppress the formation of the Ta2O5 phase. In Figure 5a, the main post-oxidation product has large elongated grains as compared to the grains in Figure 5b,c. The inset shows the top  view of the elongated grains where grains look more equiaxial. The grain enlargement is caused by the dramatic volume change associated with the formation of Ta2O5 and the very high Pilling-Bedworth ratio of 2.5. Additionally, TaC has a cubic crystal structure, whereas Ta2O5 has an orthorhombic crystal structure. The transformation from cubic to orthorhombic leads to elongation in one direction, which is also the reason why Ta2O5 has large grain size. Larger grain size is also possible due to grain growth because of the earlier onset of oxidation in the T80H20 sample as compared to the other two solid  solutions. In the oxidation of the TaC-HfC solid solutions, especially in the T80H20 sample, the formed Ta2O5 can react with the unreacted HfC. The reaction (Reaction (3)) replaces the Ta2O5 with HfO2 , so the volume change during the oxidation process is reduced. The consumption of Ta2O5 reduced the volume change and increased the mechanical integrity of (cid:16)∆G (kJ) = −0.14T − 759.76 the post-oxidation samples in solid solutions.  3Ta2O5 + 7HfC → 7HfO2 + 6TaC + CO  (cid:17)  500-2000  C  T :     [17]  (3)  The oxidation behavior of T20H80 is similar to that of pure HfC, as shown in Figure 5b. Another  beneficial effect of reaction 3 is the delay of the formation of gaseous products in the early oxidation  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa06\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa04.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0pure\\xa0TaC\\xa0and\\xa0(b)\\xa0pure\\xa0HfC.\\xa0Although\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0do\\xa0not\\xa0have\\xa0sufficient\\xa0mechanical\\xa0integrity\\xa0to\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0(Figure\\xa02c,d),\\xa0these\\xa0two\\xa0solid\\xa0solutions\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance\\xa0by\\xa0delaying\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0temperature,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0oxidation\\xa0onset\\xa0temperatures\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0are\\xa0850\\xa0and\\xa0900\\xa0°C,\\xa0respectively,\\xa0which\\xa0is\\xa0significantly\\xa0delayed\\xa0as\\xa0compared\\xa0to\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0However,\\xa0both\\xa0T80H20\\xa0and\\xa0T20H80\\xa0still\\xa0reached\\xa0near\\xa0100%\\xa0oxidation.\\xa0The\\xa0T50H50\\xa0solid\\xa0solution\\xa0displayed\\xa0the\\xa0best\\xa0oxidation\\xa0resistance.\\xa0Not\\xa0only\\xa0did\\xa0the\\xa0onset\\xa0temperature\\xa0for\\xa0oxidation\\xa0increase\\xa0to\\xa0near\\xa0940\\xa0°C\\xa0for\\xa0the\\xa0T50H50\\xa0samples,\\xa0only\\xa060%\\xa0of\\xa0the\\xa0oxidation\\xa0was\\xa0completed\\xa0when\\xa0the\\xa0temperature\\xa0reached\\xa01400\\xa0°C.\\xa0To\\xa0further\\xa0understand\\xa0the\\xa0mechanism\\xa0of\\xa0improved\\xa0oxidation\\xa0resistance\\xa0in\\xa0the\\xa0solid\\xa0solution\\xa0samples,\\xa0especially\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0the\\xa0morphologies\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0investigated\\xa0by\\xa0SEM\\xa0and\\xa0presented\\xa0in\\xa0Figure\\xa05.\\xa0As\\xa0described\\xa0earlier,\\xa0the\\xa0oxide\\xa0morphologies\\xa0of\\xa0pure\\xa0TaC\\xa0and\\xa0HfC\\xa0(Figure\\xa04)\\xa0are\\xa0highly\\xa0porous\\xa0due\\xa0to\\xa0gaseous\\xa0products\\xa0resulting\\xa0in\\xa0prominent\\xa0volume\\xa0change.\\xa0The\\xa0key\\xa0factor\\xa0in\\xa0improving\\xa0the\\xa0oxidation\\xa0resistance\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0is\\xa0to\\xa0suppress\\xa0the\\xa0formation\\xa0of\\xa0the\\xa0Ta2O5\\xa0phase.\\xa0In\\xa0Figure\\xa05a,\\xa0the\\xa0main\\xa0post‐oxidation\\xa0product\\xa0has\\xa0large\\xa0elongated\\xa0grains\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0grains\\xa0in\\xa0Figure\\xa05b,c.\\xa0The\\xa0inset\\xa0shows\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains\\xa0where\\xa0grains\\xa0look\\xa0more\\xa0equiaxial.\\xa0The\\xa0grain\\xa0enlargement\\xa0is\\xa0caused\\xa0by\\xa0the\\xa0dramatic\\xa0volume\\xa0change\\xa0associated\\xa0with\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0and\\xa0the\\xa0very\\xa0high\\xa0Pilling‐Bedworth\\xa0ratio\\xa0of\\xa02.5.\\xa0Additionally,\\xa0TaC\\xa0has\\xa0a\\xa0cubic\\xa0crystal\\xa0structure,\\xa0whereas\\xa0Ta2O5\\xa0has\\xa0an\\xa0orthorhombic\\xa0crystal\\xa0structure.\\xa0The\\xa0transformation\\xa0from\\xa0cubic\\xa0to\\xa0orthorhombic\\xa0leads\\xa0to\\xa0elongation\\xa0in\\xa0one\\xa0direction,\\xa0which\\xa0is\\xa0also\\xa0the\\xa0reason\\xa0why\\xa0Ta2O5\\xa0has\\xa0large\\xa0grain\\xa0size.\\xa0Larger\\xa0grain\\xa0size\\xa0is\\xa0also\\xa0possible\\xa0due\\xa0to\\xa0grain\\xa0growth\\xa0because\\xa0of\\xa0the\\xa0earlier\\xa0onset\\xa0of\\xa0oxidation\\xa0in\\xa0the\\xa0T80H20\\xa0sample\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0two\\xa0solid\\xa0solutions.\\xa0In\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0especially\\xa0in\\xa0the\\xa0T80H20\\xa0sample,\\xa0the\\xa0formed\\xa0Ta2O5\\xa0can\\xa0react\\xa0with\\xa0the\\xa0unreacted\\xa0HfC.\\xa0The\\xa0reaction\\xa0(Reaction\\xa0(3))\\xa0replaces\\xa0the\\xa0Ta2O5\\xa0with\\xa0HfO2,\\xa0so\\xa0the\\xa0volume\\xa0change\\xa0during\\xa0the\\xa0oxidation\\xa0process\\xa0is\\xa0reduced.\\xa0\\xa0Figure\\xa05.\\xa0Cont.\\xa0\\x0c', 'Coatings 2017, 7, 111  7 of 9  stage. Reaction 2 shows that one mole of HfC would react with one mole of oxygen to form one mole of gaseous products (CO or CO2 ). However, in reaction 3, Ta2O5 would consume seven moles of HfC to generate one mole of gaseous product, so the gases produced by HfC can be reduced or at least  delayed. The Gibbs free energy for reaction 3 is computed using Factsage [17]. The delay of the gaseous  product diminishes the cracking within the oxides and enhances the mechanical integrity of the oxide  scales in solid solutions. The well-adhered oxide scale can serve as a protection layer against the  further oxidation of solid solutions. The morphology of the post-oxidation sample of T50H50 is a vivid  proof of this concept, as shown in Figure 5c. The oxide scale of the oxidized T50H50 solid solution is  much denser as compared to the other samples. The oxide grains are mostly equiaxed, suggesting  the moderate volume increase. No large cracks are noticeable in the oxidized T50H50 sample. Thus,  the T50H50 solid solution shows the best oxidation resistance among all the solid solutions.  Figure 5. Post-oxidation SEM images of (a) T80H20 (inset showing the top view of the elongated  grains); (b) T20H80; and (c) T50H50. (Insets are high magnification images).  It must be noted that oxidation onset temperature for T50H50 is 940  C. The onset temperature reveals the interaction between the carbides and oxygen. The study of the absorption of oxygen on  TaC and HfC (001) planes suggests that the oxygen tends to sit on the Hf-C bridge [23]. In the case  of TaC, the preferential oxygen site was on the Ta-Ta bridge. After forming a solid solution, the Ta  atoms are partially replaced by Hf atoms, so the availability of Ta-Ta is disturbed, and the oxidation of the Hf element remains unaffected. Thus, the formation of Ta2O5 is retarded. As discussed earlier, the formation of Ta2O5 is detrimental to oxidation performance, so the solid solutions are expected to have superior oxidation resistances. Among the five samples, abrupt weight increase is observed only  in the pure TaC sample. The weight change in the HfC-contained samples increases steadily, which  corresponds to the adsorption of oxygen, a unique feature of HfC. The solid solutions should have  an abrupt weight increase at a lower temperature similar to pure TaC if the oxidation of TaC has not  been retarded. The maximum delay in the onset of oxidation is found in the T50H50 sample, which is  expected as the Ta-Ta bridges have been disturbed the most by forming solid solutions.  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa06\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa04.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0pure\\xa0TaC\\xa0and\\xa0(b)\\xa0pure\\xa0HfC.\\xa0Although\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0do\\xa0not\\xa0have\\xa0sufficient\\xa0mechanical\\xa0integrity\\xa0to\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0(Figure\\xa02c,d),\\xa0these\\xa0two\\xa0solid\\xa0solutions\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance\\xa0by\\xa0delaying\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0temperature,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0oxidation\\xa0onset\\xa0temperatures\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0are\\xa0850\\xa0and\\xa0900\\xa0°C,\\xa0respectively,\\xa0which\\xa0is\\xa0significantly\\xa0delayed\\xa0as\\xa0compared\\xa0to\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0However,\\xa0both\\xa0T80H20\\xa0and\\xa0T20H80\\xa0still\\xa0reached\\xa0near\\xa0100%\\xa0oxidation.\\xa0The\\xa0T50H50\\xa0solid\\xa0solution\\xa0displayed\\xa0the\\xa0best\\xa0oxidation\\xa0resistance.\\xa0Not\\xa0only\\xa0did\\xa0the\\xa0onset\\xa0temperature\\xa0for\\xa0oxidation\\xa0increase\\xa0to\\xa0near\\xa0940\\xa0°C\\xa0for\\xa0the\\xa0T50H50\\xa0samples,\\xa0only\\xa060%\\xa0of\\xa0the\\xa0oxidation\\xa0was\\xa0completed\\xa0when\\xa0the\\xa0temperature\\xa0reached\\xa01400\\xa0°C.\\xa0To\\xa0further\\xa0understand\\xa0the\\xa0mechanism\\xa0of\\xa0improved\\xa0oxidation\\xa0resistance\\xa0in\\xa0the\\xa0solid\\xa0solution\\xa0samples,\\xa0especially\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0the\\xa0morphologies\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0investigated\\xa0by\\xa0SEM\\xa0and\\xa0presented\\xa0in\\xa0Figure\\xa05.\\xa0As\\xa0described\\xa0earlier,\\xa0the\\xa0oxide\\xa0morphologies\\xa0of\\xa0pure\\xa0TaC\\xa0and\\xa0HfC\\xa0(Figure\\xa04)\\xa0are\\xa0highly\\xa0porous\\xa0due\\xa0to\\xa0gaseous\\xa0products\\xa0resulting\\xa0in\\xa0prominent\\xa0volume\\xa0change.\\xa0The\\xa0key\\xa0factor\\xa0in\\xa0improving\\xa0the\\xa0oxidation\\xa0resistance\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0is\\xa0to\\xa0suppress\\xa0the\\xa0formation\\xa0of\\xa0the\\xa0Ta2O5\\xa0phase.\\xa0In\\xa0Figure\\xa05a,\\xa0the\\xa0main\\xa0post‐oxidation\\xa0product\\xa0has\\xa0large\\xa0elongated\\xa0grains\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0grains\\xa0in\\xa0Figure\\xa05b,c.\\xa0The\\xa0inset\\xa0shows\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains\\xa0where\\xa0grains\\xa0look\\xa0more\\xa0equiaxial.\\xa0The\\xa0grain\\xa0enlargement\\xa0is\\xa0caused\\xa0by\\xa0the\\xa0dramatic\\xa0volume\\xa0change\\xa0associated\\xa0with\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0and\\xa0the\\xa0very\\xa0high\\xa0Pilling‐Bedworth\\xa0ratio\\xa0of\\xa02.5.\\xa0Additionally,\\xa0TaC\\xa0has\\xa0a\\xa0cubic\\xa0crystal\\xa0structure,\\xa0whereas\\xa0Ta2O5\\xa0has\\xa0an\\xa0orthorhombic\\xa0crystal\\xa0structure.\\xa0The\\xa0transformation\\xa0from\\xa0cubic\\xa0to\\xa0orthorhombic\\xa0leads\\xa0to\\xa0elongation\\xa0in\\xa0one\\xa0direction,\\xa0which\\xa0is\\xa0also\\xa0the\\xa0reason\\xa0why\\xa0Ta2O5\\xa0has\\xa0large\\xa0grain\\xa0size.\\xa0Larger\\xa0grain\\xa0size\\xa0is\\xa0also\\xa0possible\\xa0due\\xa0to\\xa0grain\\xa0growth\\xa0because\\xa0of\\xa0the\\xa0earlier\\xa0onset\\xa0of\\xa0oxidation\\xa0in\\xa0the\\xa0T80H20\\xa0sample\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0two\\xa0solid\\xa0solutions.\\xa0In\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0especially\\xa0in\\xa0the\\xa0T80H20\\xa0sample,\\xa0the\\xa0formed\\xa0Ta2O5\\xa0can\\xa0react\\xa0with\\xa0the\\xa0unreacted\\xa0HfC.\\xa0The\\xa0reaction\\xa0(Reaction\\xa0(3))\\xa0replaces\\xa0the\\xa0Ta2O5\\xa0with\\xa0HfO2,\\xa0so\\xa0the\\xa0volume\\xa0change\\xa0during\\xa0the\\xa0oxidation\\xa0process\\xa0is\\xa0reduced.\\xa0\\xa0Figure\\xa05.\\xa0Cont.\\xa0Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa07\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa05.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0T80H20\\xa0(inset\\xa0showing\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains);\\xa0(b)\\xa0T20H80;\\xa0and\\xa0(c)\\xa0T50H50.\\xa0(Insets\\xa0are\\xa0high\\xa0magnification\\xa0images).\\xa0The\\xa0consumption\\xa0of\\xa0Ta2O5\\xa0reduced\\xa0the\\xa0volume\\xa0change\\xa0and\\xa0increased\\xa0the\\xa0mechanical\\xa0integrity\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0in\\xa0solid\\xa0solutions.\\xa03Ta(cid:2870)O(cid:2873)(cid:3397)7HfC→7HfO(cid:2870)(cid:3397)6TaC(cid:3397)CO\\xa0(ΔG\\xa0(kJ)\\xa0=\\xa0−0.14T\\xa0−\\xa0759.76\\xa0\\xa0\\xa0\\xa0\\xa0\\xa0T:\\xa0500-2000\\xa0°C)\\xa0[17]\\xa0(3)\\xa0The\\xa0oxidation\\xa0behavior\\xa0of\\xa0T20H80\\xa0is\\xa0similar\\xa0to\\xa0that\\xa0of\\xa0pure\\xa0HfC,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa05b.\\xa0Another\\xa0beneficial\\xa0effect\\xa0of\\xa0reaction\\xa03\\xa0is\\xa0the\\xa0delay\\xa0of\\xa0the\\xa0formation\\xa0of\\xa0gaseous\\xa0products\\xa0in\\xa0the\\xa0early\\xa0oxidation\\xa0stage.\\xa0Reaction\\xa02\\xa0shows\\xa0that\\xa0one\\xa0mole\\xa0of\\xa0HfC\\xa0would\\xa0react\\xa0with\\xa0one\\xa0mole\\xa0of\\xa0oxygen\\xa0to\\xa0form\\xa0one\\xa0mole\\xa0of\\xa0gaseous\\xa0products\\xa0(CO\\xa0or\\xa0CO2).\\xa0However,\\xa0in\\xa0reaction\\xa03,\\xa0Ta2O5\\xa0would\\xa0consume\\xa0seven\\xa0moles\\xa0of\\xa0HfC\\xa0to\\xa0generate\\xa0one\\xa0mole\\xa0of\\xa0gaseous\\xa0product,\\xa0so\\xa0the\\xa0gases\\xa0produced\\xa0by\\xa0HfC\\xa0can\\xa0be\\xa0reduced\\xa0or\\xa0at\\xa0least\\xa0delayed.\\xa0The\\xa0Gibbs\\xa0free\\xa0energy\\xa0for\\xa0reaction\\xa03\\xa0is\\xa0computed\\xa0using\\xa0Factsage\\xa0[17].\\xa0The\\xa0delay\\xa0of\\xa0the\\xa0gaseous\\xa0product\\xa0diminishes\\xa0the\\xa0cracking\\xa0within\\xa0the\\xa0oxides\\xa0and\\xa0enhances\\xa0the\\xa0mechanical\\xa0integrity\\xa0of\\xa0the\\xa0oxide\\xa0scales\\xa0in\\xa0solid\\xa0solutions.\\xa0The\\xa0well‐adhered\\xa0oxide\\xa0scale\\xa0can\\xa0serve\\xa0as\\xa0a\\xa0protection\\xa0layer\\xa0against\\xa0the\\xa0further\\xa0oxidation\\xa0of\\xa0solid\\xa0solutions.\\xa0The\\xa0morphology\\xa0of\\xa0the\\xa0post‐oxidation\\xa0sample\\xa0of\\xa0T50H50\\xa0is\\xa0a\\xa0vivid\\xa0proof\\xa0of\\xa0this\\xa0concept,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa05c.\\xa0The\\xa0oxide\\xa0scale\\xa0of\\xa0the\\xa0oxidized\\xa0T50H50\\xa0solid\\xa0solution\\xa0is\\xa0much\\xa0denser\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0samples.\\xa0The\\xa0oxide\\xa0grains\\xa0are\\xa0mostly\\xa0equiaxed,\\xa0suggesting\\xa0the\\xa0moderate\\xa0volume\\xa0increase.\\xa0No\\xa0large\\xa0cracks\\xa0are\\xa0noticeable\\xa0in\\xa0the\\xa0oxidized\\xa0T50H50\\xa0sample.\\xa0Thus,\\xa0the\\xa0T50H50\\xa0solid\\xa0solution\\xa0shows\\xa0the\\xa0best\\xa0oxidation\\xa0resistance\\xa0among\\xa0all\\xa0the\\xa0solid\\xa0solutions.\\xa0It\\xa0must\\xa0be\\xa0noted\\xa0that\\xa0oxidation\\xa0onset\\xa0temperature\\xa0for\\xa0T50H50\\xa0is\\xa0940\\xa0°C.\\xa0The\\xa0onset\\xa0temperature\\xa0reveals\\xa0the\\xa0interaction\\xa0between\\xa0the\\xa0carbides\\xa0and\\xa0oxygen.\\xa0The\\xa0study\\xa0of\\xa0the\\xa0absorption\\xa0of\\xa0oxygen\\xa0on\\xa0TaC\\xa0and\\xa0HfC\\xa0(001)\\xa0planes\\xa0suggests\\xa0that\\xa0the\\xa0oxygen\\xa0tends\\xa0to\\xa0sit\\xa0on\\xa0the\\xa0Hf-C\\xa0bridge\\xa0[23].\\xa0In\\xa0the\\xa0case\\xa0of\\xa0TaC,\\xa0the\\xa0preferential\\xa0oxygen\\xa0site\\xa0was\\xa0on\\xa0the\\xa0Ta-Ta\\xa0bridge.\\xa0After\\xa0forming\\xa0a\\xa0solid\\xa0solution,\\xa0the\\xa0Ta\\xa0atoms\\xa0are\\xa0partially\\xa0replaced\\xa0by\\xa0Hf\\xa0atoms,\\xa0so\\xa0the\\xa0availability\\xa0of\\xa0Ta-Ta\\xa0is\\xa0disturbed,\\xa0and\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0Hf\\xa0element\\xa0remains\\xa0unaffected.\\xa0Thus,\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0is\\xa0retarded.\\xa0As\\xa0discussed\\xa0earlier,\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0is\\xa0detrimental\\xa0to\\xa0oxidation\\xa0performance,\\xa0so\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0expected\\xa0to\\xa0have\\xa0superior\\xa0oxidation\\xa0resistances.\\xa0Among\\xa0the\\xa0five\\xa0samples,\\xa0abrupt\\xa0weight\\xa0increase\\xa0is\\xa0observed\\xa0only\\xa0in\\xa0the\\xa0pure\\xa0TaC\\xa0sample.\\xa0The\\xa0weight\\xa0change\\xa0in\\xa0the\\xa0HfC‐contained\\xa0samples\\xa0increases\\xa0steadily,\\xa0which\\xa0corresponds\\xa0to\\xa0the\\xa0adsorption\\xa0of\\xa0oxygen,\\xa0a\\xa0unique\\xa0feature\\xa0of\\xa0HfC.\\xa0The\\xa0solid\\xa0solutions\\xa0should\\xa0have\\xa0an\\xa0abrupt\\xa0weight\\xa0increase\\xa0at\\xa0a\\xa0lower\\xa0temperature\\xa0similar\\xa0to\\xa0pure\\xa0TaC\\xa0if\\xa0the\\xa0oxidation\\xa0of\\xa0TaC\\xa0has\\xa0not\\xa0been\\xa0retarded.\\xa0The\\xa0maximum\\xa0delay\\xa0in\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0is\\xa0found\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0which\\xa0is\\xa0expected\\xa0as\\xa0the\\xa0Ta-Ta\\xa0bridges\\xa0have\\xa0been\\xa0disturbed\\xa0the\\xa0most\\xa0by\\xa0forming\\xa0solid\\xa0solutions.\\xa0\\xa0\\xa0\\x0c', 'Coatings 2017, 7, 111  4. Conclusions  8 of 9  Through TGA analyses, we investigated the oxidation behavior of pure TaC and HfC as well as  their solid solutions. The solid solutions exhibit improved oxidation resistance compared to the pure  carbides. T50H50 is found to have the best oxidation resistance, followed by T20H80 and T80H20. The onset of oxidation in T50H50 increases by 170 and 120  C as compared to pure TaC and pure HfC, respectively. The improved oxidation resistance can be attributed to the formation of the solid solutions that disturbs the atomic arrangement. Such disturbance delays the formation of Ta2O5 and does not affect the formation of HfO2 . The reaction between Ta2O5 and HfC is also responsible for the superior oxidation performances in the solid solution samples. It diminishes the generation of  gaseous products during oxidation, which reduces the porosity of the oxide scales and leads to the  better protection of the underlying materials. The present study showcases SPS-sintered solid solutions  as a new class of oxidation-resistant materials within ultrahigh temperature ceramics (UHTCs).  Acknowledgments: Cheng Zhang thanks the Florida International University Graduate School for the Dissertation Year Fellowship (DYF) award. Advanced Materials Engineering Research Institute (AMERI), FIU is acknowledged for the research facilities used and the support from its staff in this study.  Author Contributions: Benjamin Boesl and Arvind Agarwal conceived and designed experiments; Archana Loganathan performed the experiments; Cheng Zhang, Archana Loganathan, Benjamin Boesl and Arvind Agarwal analyzed data; Cheng Zhang wrote the paper.  Conflicts of Interest: The authors declare no conflict of interest.  References  1.  2.  3.  4.  5.  Fahrenholtz, W.G.; Wuchina, E.J.; Lee, W.E.; Zhou, Y. Ultra-High Temperature Ceramics: Materials for Extreme  Environment Applications; John Wiley & Sons, Inc.: Hoboken, NJ, USA, 2014.  Louis, L.E. Transition Metal Carbides and Nitrides; Academic Press: New York, NY, USA, 1971.  Upadhya, K.; Yang,  J.; Hoffman, W.P. Materials  for Ultrahigh Temperature Structural Applications.  Am. Ceram. Soc. Bull. 1997, 76, 51-56.  Pierson, H.O. Handbook of Refractory Carbides and Nitrides; William Andrew Publishing: Westwood, NJ, USA,  1996. Opeka, M.M.; Talmy, I.G.; Zaykoski, J.A. Oxidation-based Materials Selection for 2000  C + Hypersonic 2004, 39, 5887-5904.  Aerosurfaces: Theoretical Considerations and Historical Experience.  J. Mater. Sci.  [CrossRef]  6.  Simonenko, E.P.; Sevast’yanov, D.V.; Simonenko, N.P.; Sevast’yanov, V.G.; Kuznetsov, N.T. Promising  Ultra-high Temperature Ceramic Materials for Aerospace Applications. Russ.  J. Inorg. Chem.  2013, 58,  7.  8.  9.  1669-1693. [CrossRef]  Gary, S.P.; Krishnamurthy, N.; Awasthi, A.; Venkatraman, M. The O-Ta (Oxygen-Tantalum) System. J. Phase  Equilib. 1996, 17, 63-77.  Fahrenholtz, W.G.; Hilmas, G.E. Oxidation of Ultra-high Temperature Transition Metal Diboride Ceramics.  Int. Mater. Rev. 2012, 57, 61-72. [CrossRef]  Gasch, M.; Ellerby, D.; Irby, E.; Beckman, S.; Gusman, M.; Johnson, S. Processing, Properties and Arc Jet  Oxidation of Hafnium Diboride/Silicon Carbide Ultra High Temperature Ceramics. J. Mater. Sci. 2004, 39,  5925-5937. [CrossRef]  10.  Agte, C.; Alterhum, H. Investigations of the High-Melting Carbide Systems Connected with Problem of the  Carbon Melting. Z. Technol. Phyzik 1930, 11, 182-191.  11.  Coutright, E.L.; Prater, J.T.; Holcomb, G.R.; Stpierre, G.R.; Rapp, R.A. Oxidation of Hafnium Carbide and  Hafnium Carbide with Additions of Tantalum and Praseodymium. Oxid. Met. 1991, 36, 423-437. [CrossRef]  12.  13.  Ghaffari, S.A.; Faghihi-Sani, M.A.; Golestani-Fard, F.; Ebrahimi, S. Pressureless Sintering of Ta0.8Hf0.2C UHTC in the Presence of MoSi2 . Ceram. Int. 2013, 39, 1985-1989. [CrossRef] Patterson, M.C.L. Advanced HfC-TaC Oxidation Resistance Composite Rocket Thruster. Mater. Manuf.  Process. 1996, 11, 367-379. [CrossRef]  \\x0c', 'Coatings 2017, 7, 111  9 of 9  14.  Rudy, E. Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon System. Part II. Ternary Systems, Vol I.  Ta-HfC-C System; Technical Report: AFML-TR-65-2 Part II Vol. 1; Wright-Patterson Air Force Base, Air Force  Systems Command, Air Force Materials Laboratory: Dayton, OH, USA, 1969.  15.  Ghaffari, S.A.; Faghihi-Sani, M.A.; Golestani-Fard, F.; Nojabayy, M. Diffusion and Solid Solution Formation  Between the Binary Carbides of TaC, HfC, ZrC. Int. J. Refract. Met. Hard Mater. 2013, 41, 180-184. [CrossRef]  16.  Cedillos-Barraza, O.; Grasso, S.; Nasiri, N.A.; Jayaseelan, D.D.; Reece, M.J.; Lee, W.E. Sintering Behavior,  Solid Solution Formation and Characterization of TaC, HfC and TaC-HfC Fabricated by Spark Plasma  Sintering. J. Eur. Ceram. Soc. 2016, 36, 1539-1548. [CrossRef]  17.  Zhang, C.; Gupta, A.; Seal, S.; Boesl, B.; Agarwal, A. Solid Solution Synthesis of Tantalum Carbide-Hafnium  Carbide by Spark Plasma Sintering. J. Am. Ceram. Soc. 2017, 100, 1773-2308. [CrossRef]  18.  Zhang, C. High Temperature Oxidation Study of Tantalum Carbide-Hafnium Carbide Solid Solutions  Synthesized By Spark Plasma Sintering. Ph.D. Thesis, Florida International University, Miami, FL, USA,  18 October 2016.  19.  Desmaison-Brut, M.; Alexandre, N.; Desmaison, J. Comparison of the Oxidation Behavior of Two Dense Hot  Isostatically Pressed Tantalum Carbide (TaC and Ta2C) Materials. [CrossRef]  J. Eur. Ceram. Soc. 1997, 17, 1325-1334.  20.  Zhang, X.; Hilmas, G.E.; Fahrenholtz, W.G. Densification, Mehcanical Properties, and Oxidation Reistance of  TaC-TaB2 Creamics. J. Am. Ceram. Soc. 2008, 91, 4129-4132.  21.  Cramer, S.D.; Covino, B.S., Jr. ASM Handbook Volume 13A: Corrosion: Fundamentals, Testing, and Protection;  ASM International: Geauga County, OH, USA, 2013.  22.  Bargeron, C.B.; Benson, R.C.; Jette, A.N.; Phillips, T.E. Oxidation of Hafnium Carbide in the Temperature Range 1400  to 2060  C. J. Am. Ceram. Soc. 1993, 76, 1040-1046. [CrossRef]  23.  Liu, D.; Deng, J.; Jin, Y.; He, C. Adsorption of Atomic Oxygen on HfC and TaC (110) Surface From Firtst  Principles. Appl. Surf. Sci. 2012, 261, 214-218. [CrossRef]  © 2017 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access  article distributed under the terms and conditions of the Creative Commons Attribution  (CC BY) license (http://creativecommons.org/licenses/by/4.0/).  \\x0c']"
},{
  "_id": 29,
  "PDF": "Combined effects of WC and SiC on densification and thermo-mechanical stability of ZrB 2 ceramics.pdf",
  "Text": "['Materials and Design 109 (2016) 396-407  Contents lists available at ScienceDirect  Materials and Design  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t d e s  Combined effects of WC and SiC on densification and thermo-mechanical stability of ZrB2 ceramics  Frédéric Monteverde, Laura Silvestroni ⁎  CNR-ISTEC, National Research Council of Italy - Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, RA, Italy  H I G H L I G H T S  G R A P H I C A L  A B S T R A C T  • Fabrication of ZrB2-WC-SiC composites with density above 94%. • Development of “core-rim” substructures in the matrix and mixed W, Zr- boride and carbide. • Room temperature strength passed from 540 to 630 MPa to over 700 MPa at 1500 °C in air. • W, enclosed in the solid solution shells, retarded the oxygen diffusion compared to pure ZrB2.  a r t i c l e  i n f o  a b s t r a c t  ZrB2-based mixtures containing WC and SiC in various amounts were hot-pressed at 1930 °C, achieving final relative densities above 94%. The diboride matrix of the sintered ceramics was constituted by (Zr,W)B2 solid solution shells grown onto original ZrB2 grains. WC phase reacted during sintering leaving W-monoboride and mixed Zr,W-carbide as secondary phases. Room temperature flexure strength ranged from 540 to 630 MPa, but performances of major significance were determined at 1500 °C in air, with values exceeding 700 MPa. SiC was vital to improve the oxidation resistance at 1500 °C compared to SiC-free ZrB2-based ceramic. W, encased in the shell, retarded the inward diffusion of oxygen compared to the pure ZrB2. © 2016 Published by Elsevier Ltd.  Article history:  Received 13 April 2016 Received in revised form 24 June 2016 Accepted 28 June 2016 Available online 1 July 2016  Keywords:  ZrB2 Tungsten Solid solution Mechanical properties Oxidation  1. Introduction  Zirconium diboride (ZrB2) belongs to the well-known class of compounds conventionally named as ultra-high-temperature ceramics (UHTCs) because of their melting point above 3000 °C. Research activity of recent years has been addressed to continuously improve its thermo ⁎  Corresponding author. E-mail address: laura.silvestroni@istec.cn.rit (L. Silvestroni).  http://dx.doi.org/10.1016/j.matdes.2016.06.114 0264-1275/© 2016 Published by Elsevier Ltd.  mechanical capabilities, failure tolerance, and ablation/oxidation resistance through the tailored addition of various secondary phases [1-5]. SiC has been repeatedly added in a range of 5-30 vol% to transition metal diborides (MeB2), owing to its proven ability to help densification of MeB2, as well as to limit the growth rate of the diboride matrix by pinning the grain boundary migration during the final stage of densification, typically above 1900 °C in the case of hot-pressing [1,6]. The presence of SiC has also proved to increase flexure strength owing to a refined and dense microstructure [7], as well as the resistance to  \\x0c', \"F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  397  oxidation up to 1650 °C, thanks to the in situ formation of a multilayered protective oxide scale [8,9]. Additional third phases able to bring benefit to the high temperature performance of MeB2-SiC systems include carbides and silicides of W, Ta, or Mo [3,10-13]. In this respect, the incorporation of WC has recently gained great attention thanks to an attractive ability [10,14,15] to favor the flexure strength retention at temperatures up to 1600 °C. Also the resistance to oxidation was reported to take advantages from the presence of W-carrying refractory phases inside the resulting microstructure of ZrB2 ceramics [16-19]. No matter if WC is initially added as second phase [19] or as impurity [7], improved performances at high temperatures were recorded anyways. Average flexure strengths up to 640 MPa at 1600 °C in protective Ar atmosphere were achieved as long as a starting quantity of 5 vol% WC was introduced into a ZrB2-20 vol% SiC powder mix [10]. The potential of the ZrB2- SiC-WC system has been further disclosed thanks to measurements of residual flexure strength (tested at room temperature) after 24 h of exposure to static air at 1400 °C [20]. This material, object of the present work and thereafter designated as ZS15-5WC, despite the adverse influence of 6 vol% residual porosity, lost only 28% of its average room temperature flexure strength, decreasing from 543 to 390 MPa. As far as the resistance to oxidation is concerned, the beneficial effects to the oxidation rates of ZrB2 and ZrB2-SiC ceramics which arise from the presence of tungsten has been already outlined in previous works [10,16-18,21]. The improvement was ascribed to a number of phenomena: i) the occurrence of an eutectic between WO3 and ZrO2 at 1275 °C activates a liquid phase sintering of the in situ growing zirconia scale decreasing its porosity [10,16]; ii) the oxidation product, WO3, performs as barrier itself due to the volume increase associated to oxidation of W-species to WO3 [16]; iii) the incorporation of W atoms in the borate glass increases the stability of the glass itself [17,18]. In this study, we report on advances in the understanding of the densification and oxidation behavior of W-doped ZrB2-based ceramics and their thermo-mechanical improvements attained through the addition of refractory carbides such as SiC and WC. The main novelty of this flexure work is related to the achievement of strength values overpassing 700 MPa when tested at 1500 °C in air.  2. Experimental procedure  Commercial powders of ZrB2, SiC, and WC (see Table 1) were used to produce three different compositions in the ZrB2-WC-SiC system. The initial nominal compositions were all based on ZrB2, with SiC and/or WC as secondary phases in varying amounts, as follows (vol%):  ZS0-15WC: ZS3-5WC: ZS15-5WC:  ZrB2 + 15WC ZrB2 + 3SiC + 5WC ZrB2 + 15SiC + 5WC  With these three compositions, multiple purposes were addressed. ZS0-15WC was designed to keep the system as simple as possible to study the reaction products during sintering. In ZS15-5WC, the amount of SiC was set as that typically introduced in the majority of the studies on ZrB2-SiC [1-3], while ZS3-5WC was conceived as optimization of the previous two compositions, in an effort to reduce the amount of SiC phase under the percolation limit, as main responsible of strength decrease at high temperature.  Base component (ZrB2) and secondary phases (SiC, WC) were weighed according to the rule-of-mixture, batched into a PET bottle and then ball-milled for 24 h in absolute ethanol using SiC milling media. Subsequently, the slurries, dried using a rotary evaporator, were sieved through a 150 μm metallic screen. 30 mm diameter pellets were cold compacted using a uniaxial press with 20 MPa applied pressure. The green pellets were directly positioned into the graphite die located inside the vacuum chamber of the hot-press furnace. The inner walls of the graphite die were precautionary protected by a BNsprayed graphitized 0.5 mm thick foil. The pellets were thus hotpressed in the vacuum range of 0.1-1 mbar using an induction-heated graphite die and applying a constant uniaxial pressure of 30 MPa during ramping up to 1930 °C, 15-20 °C/min of heating rate. Once the target temperature was reached, the linear pressure was increased up to 40 MPa and maintained up to 35 min to promote full densification. The change in thickness (dL) of the green pellets under hot-pressing was recorded during heating up and final dwell. The experimental dL values were then converted into relative density data, under the hypothesis of a full mass conservation. The temperature was monitored by means of an optical pyrometer focused into a blind hole dug on the external surface of the graphite die. At the end of the dwell, free cooling followed. Crystalline phases of the sintered ceramics were identified by X-ray diffraction (XRD, mod. D8 Advance - Bruker, Germany). The cell parameters of the secondary phases indexed in the hot-pressed ZrB2 ceramics were determined using the Rietveld refinement approach. Specific XRD patterns were acquired on a polished surface of a bulk sample, imposing the strongest tabulated ZrB2 reflections along the 2-theta scanned range of 20°-80° as the internal calibrating known phase. The microstructure of the as-hot-pressed ZrB2 ceramics were also analyzed on fractured and polished surfaces by field-emission scanning electron microscopy (FESEM, mod. ΣIGMA, ZEISS NTS Gmbh, Germany) coupled to an energy dispersive X-ray micro-analyzer (EDS, mod. INCA Energy 300, Oxford instruments, UK). Key microstructural features like residual porosity, mean grain size, and volumetric content of the secondary phases were evaluated thanks to FESEM micrographs elaborated with the support of the commercial software package Image Pro Plus (v.7, Media Cybernetics, USA). The bulk densities (ρB) of the as-hot-pressed ceramics were measured by Archimedes' method in distillate water at 25 °C, while the theoretical densities (ρTH,0) were computed according to the rule-of-phase mixture. Residual porosity was first calculated dividing ρB by ρTH,0 and then adjusted/confirmed by scanning electron microscopy inspection on polished sections. The expected final densities (ρTH,1) were updated according to the estimated quantities of the secondary phases. The final relative densities (RD*) were established according to the image analysis approach. The 4-pt. flexure strength (σ) was measured at room temperature (RT) and at 1500 °C in air using the guidelines of the European standards for advanced ceramics ENV843-1:2006 and EN820-1:2002, respectively. Chamfered bars with dimensions 25.0 × 2.5 × 2.0 mm3 (length by width by thickness, respectively) were tested at RT using a semiarticulated steel-made 4-pt fixture (lower span 20 mm, upper span 10 mm) in a screw-driven load frame (Zwick-Roell mod. Z050, Germany), 1 mm/min of cross-head speed. Flexure strength at high temperature was instead measured using an Instron apparatus (mod. 6025) equipped with a 4-pt fixture made of SiC. Before applying the load during testing at 1500 °C, a dwell of 18 min was set to reach  Table 1 Main characteristics of the starting powders used. ρTH: theoretical density, d50: median of the granulometric distribution.  Powder  ZrB2 SiC WC  Grade and supplier  ρTH (g/cm3)  d50 (μm)  Purity (wt%)  Main impurities (wt%)  Gr. B, H.C. Starck, Germany Gr. BF-12, H.C. Starck, Germany Gr. SD0.5, Treibacher Industries AG, Germany  6.1 3.21 15.74  2.1 0.6 0.5  N 98 N 99 99.5  O:1.5, Hf:0.2 O:0.9  -  \\x0c\", '398  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  thermal equilibrium. For each temperature, at least 3 bars per composition were tested. The resistance to oxidation was studied using a bottom-loading furnace (Nannetti FC18, Faenza, Italy) exposing rectangular coupons 13.0 × 2.5 × 2 mm3 for 15 min to the effects of a stagnant air at 1500 °C. The dwell time was chosen on the basis of typical durations of re-entry phase of space shuttles. The coupons, formerly cleaned in acetone and weighed, were bottom-up loaded into the hot zone of the furnace when the target temperature was already achieved, elapsed 15 min, and then rapidly down-loaded to allow air-quenching. Postoxidation changes were first evaluated by calculating the specific mass variation, i.e. mass change (Δm) over the exposed initial area (A). The microstructural modifications induced by the oxidation attack were analyzed by XRD and FESEM-EDS analyses on the external oxidized surfaces and fractured cross-sections.  3. Results and discussion  3.1. X-ray diffraction analysis  The X-ray diffraction patterns of the three hot-pressed ZrB2 ceramics in Fig. 1 show the hexagonal ZrB2 as the dominant phase (hex-ZrB2, PDF#34-0423). In addition, two other non-stoichiometric phases based on tetragonal tungsten boride (t-WB, PDF#35-0738) and cubic zirconium carbide (c-ZrC, PDF#35-0784) were also indexed, in larger amount for the ZS0-15WC ceramic. Based on the peak intensity ratio, the amount of the tetragonal W-boride does not differ too much in the ZS3-5WC and ZS15-5WC composites, coherently with the same starting WC amount of 5 vol%. On the other hand, residual WC fell below the XRD detection limit and was not indexed in any ceramics object of the present work. The refined cell parameters of the secondary phases indexed as t-WB and c-ZrC provided evidence of non-negligible deviations from the expected values of the pure standard compounds, as their peaks position shifted to higher 2-theta angles (see Table 2). In accordance with other authors [13,22-25], two solid solutions based on cubic (Zr, W) carbide and tetragonal (W, Zr) boride were supposed to form upon sintering. These secondary phases were interpreted as the products of chemical reactions occurred between the starting secondary phase WC with ZrB2 particles and the oxygen-bearing species covering their external surface. The presence of a foreign element like W hosted into the cubic ZrC lattice was here confirmed by FESEM-EDS analyses, as commented later. In the case of the tetragonal WB phase, due to very scarce amount of Zr involved, together with an unfavorable overlapping of the characteristic X-ray signals generated by W and Zr, the capture of Zr into the t-WB lattice was considered sound and very probable, but in the need of an ultimate quantitative determination. The major factor affecting the overall volume change of the c(Zr,W)C and t-(W,Zr)B lattices was attributed to the substitution of W or Zr atoms, 1.42 Å and 1.59 Å of covalent radius respectively, in the respective ideal lattices. As for the c-(Zr,W)C phase is concerned, the replacement of Zr atoms with smaller W ones led the cubic lattice to shrink, Table 2. In the case of the t-(W,Zr)B phase, an anisotropic variation of the cell parameters took place, with a more pronounced expansion of the c-axis compared to the overall contraction of the basal planes, due to the new positioning of covalently larger Zr atoms, Table 2. However, the net effect for the tetragonal cell volume was surprisingly just a slight contraction in the range of 0.2-0.3%. The elemental analysis by FESEM-EDS of the t-(W,Zr)B and c-(Zr,W)C phases is more widely explored in Section 3.2. It should also be considered that the refined cell parameters shown in Table 2 may be affected to some extent by the residual stresses built-up in the sintered bulks upon cooling. The mismatch of the thermal contraction between the primary diboride matrix (7.0 ppm/K [26]) and the secondary phases (3 ppm/K for SiC [27], 7.2 ppm/K for ZrC [27], 6.5 ppm/K for WB [28]) may lead the crystal lattices of all the  Fig. 1. X-ray diffraction patterns of the hot-pressed ZrB2-ceramics: a) ZS0-15WC, b) ZS35WC, and c) ZS15-5WC.  solid phases to deform: the associated stress state gives rise to measurable shifts of the characteristic planar spacing in the XRD spectra. In fact, taking as example the ZS15-5WC ceramic, where the expected strongest reflections of β-SiC are better defined with no overlapping, the cell parameter of the cubic SiC lattice shrinks of about 0.6% compared to a strain-free state. Thankfully, the thermal expansion mismatch between the diboride matrix and c-(Zr,W)C or t-(W,Zr)B does not differ  \\x0c', 'F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  399  Table 2 Nominal and refined cell parameters ± 1 estimated standard deviation (esd) of the indexed secondary phases through the Rietveld processing approach. Vi: cell volume of the refined structures, dVi = (Vi − V0)/V0: cell volume variation with V0 cell volume of tabulated t-WB and c-ZrC.  Label  Nominal cell parameters  Refined cell parameters   Cell volume  t-WB  c-ZrC  t-(W,Zr)B  c-(Zr,W)C  t-(W,Zr)B  c-(Zr,W)C  PDF#35-0738  PDF#35-0784  a0, c0 (Å)  a0 (Å)  3.11655, 16.9101  4.6930  ZS15-5WC  ZS3-5WC  ZS0-15WC  a1 (Å)  3.11034  3.11115  3.11152  19  20  32  c1 (Å)  16.923  16.928  16.919  20  14  21  a1 (Å)  4.6327  4.6228  3  4  4.6270  12  V1 (Å3)  163.72  163.88  163.81  dV1 (%)  -0.32  -0.23  -0.27  V2 (Å3)  99.43  98.79  99.06  dV2 (%)  -3.80  -4.42  -4.16  appreciably, thus the calculated shifts of the cell volume, as shown in Table 2, can be fully attributed to composition variation of the stoichiometric t-WB and c-ZrC phases.  3.2. Microstructure analysis by FESEM-EDS  A feature common to all the as-sintered ZrB2 ceramics is the distinct morphology of the diboride matrix. ZrB2 grains developed a round and regular shape associated to a substructure, widely known as “corerim.” Such substructure consists of stoichiometric ZrB2 nuclei, i.e. cores, surrounded by a isostructural (Zr,W)B2 solid solution which constitutes the rims, as exemplified in Fig. 2. The “core-rim” feature is well developed throughout the bulk, much more pronounced in ZS0-15WC material for which the peculiar feature of the rim tends to take a continuous network likeness. The amount of W captured into the rims was investigated by FESEM-EDS applying low voltage to the primary electron beam in order to confine the excited X-ray generation volume just underneath the rim under analysis. Several FESEM-EDS measurements revealed a W content in the (Zr,W)B2 solid solution between 2 and 4 at.% with respect to Zr. Concerning the new secondary phases formed, FESEM-EDS analyses could only confirm representative compositions ranges of (W1− xZrx)B and (Zr1− yWy)C, where 0.02 ≤ x ≤ 0.04, 0.13 ≤ y ≤ 0.15. The actual expected density (ρTH,1) of the hot-pressed ceramics were then refined using the following theoretical density for t-(W,Zr)B and c-(Zr,W)C, 15.52 ± 0.08 and 7.50 ± 0.06 g/cm3, respectively, in conjunction to their volumetric estimates obtained via image analysis. The theoretical density of the aforementioned secondary phases was calculated applying the elemental composition estimated by FESEM-EDS to the refined cell volume in Table 2 obtained via the Rietveld approach. These findings, together with the formation of secondary phases, like (Zr,W)C and (W,Zr)B, demonstrate that WC is not stable across the processing temperature range, but actively interacted not only with the ZrB2 matrix, but also with its native oxide impurities, absent indeed in the final microstructure. Similar conclusions were drawn for other sintering additives used for ZrB2 or HfB2, like the disilicides of Zr,  Hf, Ta, Mo, and W [3,19], or other W-based compounds [23]. The assessment of the regular shape of the diboride matrix is of great value, because it can be linked to specific mechanisms of mass transport during densification, as discussed later on. FESEM-EDS analyses did not allow to assess in a robust manner the eventual capture of minor content of B or C into the lattice of c-(Zr,W)C or t-(W,Zr)B phases, respectively, due to the in-situ electron beam activated carbon contamination together with a strong overlapping of the excited characteristic peaks of interest for the desired analysis.  3.2.1. The ZS0-15WC ceramic  The microstructure of this ZrB2 ceramic, depicted in Fig. 3a, shows that the different phases are homogeneously distributed and no large defects can be seen. The mean grain size of the diboride matrix is 1.6 μm and 19.5 vol% W-containing secondary phases were estimated, thanks to a brighter appearance compared to the grey background associated to the diboride-based matrix. Based on the image analysis estimates, 13.9 vol% (characterized by a white smooth appearance in Fig. 3b) was attributed to (W,Zr)B, while 5.8 vol% (featured with a greyer contrast in Fig. 3b) to (Zr,W)C. It is essential repeating that these secondary phases in such amounts, as well as the extensive formation of rims surrounding ZrB2 cores cannot merely be explained in term of reactions occurring during hot-pressing between WC and the oxygen-bearing species of ZrB2, but ought to be illustrated through a direct role of the ZrB2 matrix itself. Only in the present case, due to the well-developed distinction between cores and rims, it was possible to estimate a core-to-rim volumetric ratio of 70-30 vol%. Very rarely WC with oxygen traces, distinguishable thanks to a rough bright appearance, was also found. In addition, 3 vol% residual free C is present scattered throughout the bulk, with a dark and elongated shape (see Fig. 3b). The C residue, as well as (W,Zr)B and (Zr,W)C, contributes to compose the bulk density and to make the RD value in Table 3 underestimated. According to the volumetric amount of primary and secondary phases in the hot-pressed ceramic, the re-calculation of the expected pore-less density, ρTH,1 = 7.37 ± 0.02 g/cm3 does not differ appreciably from the measured one, 7.11 ± 0.01 g/cm3. The ceramic  Fig. 2. FESEM-EDS analysis on the polished section of the ZS15-5WC ceramic showing the core (C)-rim (R) substructures with EDS spectra on the right, as well as SiC particulate (dark areas) and W boride (bright areas).  \\x0c', '400  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  Fig. 3. FESEM micrographs of polished sections: (a,b) ZS0-15WC, (c,d) ZS3-5WC, (e,f) ZS15-5WC.  results indeed fully compact and no fine closed porosity in measurable amount was seen during FESEM inspection.  3.2.2. The ZS3-5WC ceramic  FESEM images in Fig. 3c,d reveal a uniform and fully dense microstructure with regularly featured grains of the diboride matrix, 1.8 μm as mean size. The secondary phases are homogeneously distributed. SiC phase was estimated to occupy 2.8 vol% of the available volume: from this finding the authors are inclined to attribute a minor role to SiC in the chemical reactions occurred during hot-pressing that led the resulting microstructure to change significantly compared to the nominal one. In fact, 5 vol% initial WC evolved into new 6.3 vol% W containing phases, together with 0.6 vol% of free C. The peculiar “corerim” feature of the diboride matrix was clearly confirmed. SiC phase occasionally showed a “core-rim” substructure with the rim containing Zr traces, see the inset in Fig. 3d. The re-calculation of the expected density ρTH,1 was 6.45 ± 0.02 g/cm3. The overall role of SiC during hot-pressing is explained in the Section 3.3, see the inset in Fig. 3d.  3.2.3. The ZS15-5WC ceramic  The representative microstructure of this ceramic is shown in Fig. 3e,f. Differently from ZS0-15WC and ZS3-5WC, the ZS15-5WC did not reach full density, and about 6 vol% of residual porosity was observed via FESEM inspection of the polished surface. The mean grain  Table 3 Initial content of SiC and WC, main hot-pressing (HP) parameters, onset temperature (TON) of early measurable shrinkage, experimental bulk density ρB (± 0.01 g/cm3), theoretical density ρTH,0, and relative density (RD) via rule-of-starting mixture, expected final density ρTH,1 (± 0.02 g/cm3) and relative density (RD*) set via image analysis, and mean grain size (mgs ± 1 s.d) of diboride matrix.  Initial content  HP parameters  Label  SiC vol%  WC vol%  °C  min  MPa  ZS0-15WC ZS3-5WC ZS15-5WC  0 3 15  15 5 5  (1)Refined cell parameters in Table 2.  1930 1930 1930  26 30 12  40 40 40  TON  °C  1630 1550 1740  ρB  g/cm3  7.11 6.47 5.66  ρTH,0  g/cm3  7.55 6.50 6.14  RD  %  94.2 99.5 92.2  ρTH,1  g/cm3  7.37 6.45 6.03  RD*  %  96.5 N 99.9 93.8  Secondary phases by XRD(1)  mgs  μm  t-(W,Zr)B, c-(Zr,W)C t-(W,Zr)B, c-(Zr,W)C, β-SiC t-(W,Zr)B, c-(Zr,W)C β-SiC  1.6 ± 0.4 1.8 ± 0.6 1.5 ± 0.5  \\x0c', 'F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  401  size of the diboride matrix, 1.5 μm, was in the same range as the previous materials, as reported in Table 3, and the distribution of the secondary phases was confirmed homogeneous. Macro-defects or large agglomerates were not seen by FESEM. The image analysis expressed an estimate of 14.7 vol% for SiC and 7.5 vol% for W-containing secondary phases. Investigating the microstructure in thorough detail, the phases already identified in the previous ceramics were once more confirmed. A finer segmentation of the grey levels enabled to discriminate between (W,Zr)B and (Zr,W)C, and 3.5 and 4.0 vol% were set as best estimates. However, the residual free C remained undetermined owing to the widespread residual porosity featured in the FESEM micrograph with an equivalent black appearance, making the ultimate grey level segmentation for free C unsuccessful. Based on the previous findings in ZS0-15WC and ZS3-5WC ceramics, a guess of 1 vol% residual free C was supposed being formed also in ZS15-5WC material and hence used to recalculate ρTH,1, which resulted 6.03 ± 0.02 g/cm3.  3.3. Densification behavior  Looking at the relative density (RD) data summarized in Table 3, it might be inferred that the hot-pressed compositions achieved a relative density above 92%. Now, the theoretical density ρTH,0 was calculated according to the rule-of-mixture based on the only nominal starting composition. The microstructure analyses by XRD and FESEM-EDS described in Sections 3.1-3.2 proved in a robust manner that the post-sintering compositions changed significantly compared to the starting ones. Taking into account the estimates of the secondary phases via image analysis, the expected final densities, ρTH,1, were therefore re-calculated (see Table 3). Except for ZS15-5WC, the conditions adopted during hot-pressing conducted the ZS0-15WC and ZS3-5WC powder mixtures to achieve full density. The linear shrinkage (dL) recorded during the hot-press runs were elaborated under the assumption of the mass conservation, and converted into the relative density (RD*) vs. temperature/time plots (see Fig. 4). Under the same effect of 30 MPa applied pressure, the composition that started to shrink at lower temperature (TON) was ZS3-5WC (1550 °C) followed by ZS0-15WC (1630 °C) and ZS15-5WC, for which a measurable shrinkage of the sample thickness started taking place only at 1740 °C, as Fig. 4a shows. Considering the materials with the equivalent initial content of 5 vol% WC, ZS3-5WC, and ZS15-5WC, the corresponding TON values suggest that an addition of 15 vol% SiC particles is sufficient to hinder grain-boundary migration during the final stage of sintering, as reported in our previous study [3]. Fig. 4b clearly displays that ZS15-5WC composition had the most sluggish shrinkage rate along the final isothermal dwell, compared to the other two composites containing no or 3 vol% SiC.  The WC phase had unambiguously a dominant role connected to an enhanced temperature reactivity which led it to first react with ZrB2 and its native oxides and hence disappear in the final ceramic. In fact, it is commonly accepted that a metal diboride MeB2 powder, Me = Ti, Zr, Hf, can restore its intrinsic high-temperature sinterability only after having removed the oxygen-carrying species which, covering the external surfaces, passivate the overall chemical reactivity of the pure MeB2 powder [29]. A series of chemical reactions supposed to be responsible of the oxygen removal from the diboride particle surfaces and consequent formation of some reaction products have been proposed by several authors and are summarized in Table 4 [15,22,23,29]. The minimum temperature (TMIN) over which the Gibbs free energy (ΔG) of a reaction becomes negative was calculated in standard condition of isobaric pressure (1 bar) and in conditions closer to the inner working pressure during hot-pressing (1 mbar). Actually, the amounts of the secondary phases such as (W,Zr)B and (Zr,W)C experimentally estimated via image analysis in the hot-pressed ceramics cannot merely be explained according to reactions in Table 4: the oxygen impurities available for the reactions are not enough to justify the estimated final content of (W,Zr)B and (Zr,W)C. It must be thought another driving role for WC that interacts with the diboride matrix: such chemical interactions involve the entire powder mixture and are partially responsible of the formation of (Zr,W)B2, (W,Zr)B, (Zr,W)C, and free C. The reaction paths of the primary reactants (ZrB2, WC, ZrO2, and B2O3) were tentatively tracked by means of the thermodynamic solver HSC Chemistry® v. 6.12. Imposing as initial inputs 1 mol of main phases (ZrB2 and WC) and traces of oxides covering the diboride, 0.1 mol of B2O3 and ZrO2, it can be read in Fig. 5 that WC is not stable in the conditions imposed during hot-pressing: its amount progressively decreases by combining with B2O3 and forming the WB phase. Volatile CO(g) is released at the same time. Once the available B2O3 is completely consumed, the progressive dissociation of WC yields free W, which is therefore available during the whole sintering process and can be captured in the boride/carbide solid solutions. The reducing environment favors the formation of ZrC upon removal of ZrO2 and by following carburization of ZrB2. The amount of WB keeps increasing with a less steep rate by reacting with ZrB2 depleted of the native oxide contamination. The reaction path just presented is based on the assumption on specific quantities of only pure compounds. The authors are aware that little variation on one of these parameters can shift equilibria at which specific reactions take place, or local chemistry can trigger competing reactions. Presenting some outputs in Fig. 5, we overall aim at disclosing, through a thermodynamic approach, the chemical reactivity of this complex system during sintering which can give rise to new borides and carbides. Another aspect to consider in the process of densification is the possibility to form liquid phases that could enhance mass transfer  Fig. 4. Relative density (RD*) vs temperature (T) and vs. dwell time (t) of the composition ZS0-15WC, ZS3-5WC and ZS15-5WC during (a) ramping up to 1930 °C and (b) isothermal dwell up to 35 min.  \\x0c', \"402  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  Table 4 List of chemical reactions reported in literature; TMIN indicates the temperature at which Gibbs standard free energy (ΔG0) becomes negative, 1 bar or 1 mbar isobaric pressure.  #  A B C D E F G  Reactants  Products  3 WC + ZrO2 3 ZrB2 + 6 WC+ ZrO2 7 ZrB2 + 10 WC + 3 ZrO2 2 ZrB2 + 5 WC + 3 ZrO2 3 WC + B2O3 4 WC + B2O3 3 W2C + B2O3  ZrC + 3 W + 2 CO(g) 6 WB + 4 ZrC + 2 CO(g) 10 WB + 10 ZrC + 2 B2O3(g) 5 ZrC + 5 W + 2 B2O3(g) 2 WB + W + 3 CO(g) 2 WB + W2C + 3 CO(g) 2 WB + 4 W + 3 CO(g)  TMIN (°C) for ΔG0 ≤ 0  1 mbar  1550 1500 2170 2450 1230 1220 1225  Ref.  [22] [23] [23] [15] [29] [29] [29]  1 bar  1965 1810 2480 2500 1510 1500 1550  mechanisms. Thermodynamic studies in the 60-70's on the mutual stability of compounds in the W-B-C system already verified that the W2B phase has a very narrow region of homogeneity and is reduced to W metal by Ti, Ta, and Zr, and, in reducing environment, W2B forms WB, WC and a slight amount of carbon [30]. More recently, the release of W was experimentally verified by Zou and co-authors [23] who heat treated for 1 h in vacuum at 1650 and 1800 °C a mixture of WC and ZrO2; W2B can be seen as a by-product for the final W. Unfortunately, the eventual formation of another by-product like WB2 could not be taken into account because it is not included in the available HSC Chemistry database. Eutectic temperatures in the W-B binary system are foreseen from 1930 to 2600 °C, being the lowest between WB12 and B [31]. Nevertheless, the contemporary introduction of multiple elements like Zr, Si, C, and O into the W-B system can further decrease the temperature at which melts form [32] and promote the formation of a transient liquid phase (TLP) at much lower temperature, above TON. The term “liquid” is here borrowed for sake of clarity from an accepted terminology referable to the sintering of advanced ceramics. Now, the actual fluidity of TLP during HP upon its onset is here only matter of speculation. Much more important is its own transient nature: in fact the TLP disappeared upon cooling, leaving as its most tangible signature the core/rim substructure. In any case, such TLP is here designated as the main carrier capable of moving atoms through the grain boundaries to feed the diffusion-solution/re-precipitation mechanism active during HP and therefore the overall densification process. It is the present authors' belief that such an extensive formation of core-rim substructures throughout the sintered bulk must be interpreted claiming an extended mass  transfer during hot-pressing sustained by the TLP which led the ZS015WC and ZS3-5WC compositions to fully densify. As for the role of SiC particulate inside WC-containing ZrB2-based compositions (i.e. ZS3-5WC and ZS15-5WC), some considerations can be drawn. Based on thermodynamic calculations, SiC and WC do not react up to 1930 °C: the proof is that image analyses, within the accuracy of the method, basically gave back the initially batched quantity. Actually, the silica contamination of SiC, typical of commercial non-oxide powders (see Table 1), did react with WC and started playing a role during the earliest onset of the TLP formation. Fig. 6 displays the path through which WC and SiO2 are supposed to start interacting up to their full transformation into W and gaseous species, according to reaction (1) and it appears that W is available at temperature as low as 1130 °C.  lð Þ þ WC→W sð Þ þ SiO gð  Þ þ CO gð Þ  SiO2  ð1Þ  Increasing the content of SiC (i.e. ZS15-5WC), the yield of W, favorable for sintering, is neutralized by an overall enhanced refractoriness of the ZS15-5WC powder mixture. In the last cited composition, the driving force to eliminate porosity supported by the TLP mechanism was indeed adversely affected by the dragging action of 15 vol% SiC particles which partly pinned the boundaries of the diboride grains without leaving them to migrate away.  3.4. Mechanical properties  The experimental mechanical properties of the ZrB2 ceramics herein developed are summarized in Table 5 and then compared to published data on ZrB2-based ceramics initially batched with W-carrying compounds, in form of silicide or carbide [10,19,33]. Among the ceramics object of the present study, ZS15-5WC exhibited the lowest room temperature average flexure strength, 543 MPa, primarily owing to 6 vol% residual porosity, while the fully dense ZS3-5WC and ZS0-15WC ceramics exhibited a statistically equivalent mean value, around 630 MPa. When the SiC-free ceramic ZS0-15WC was tested at 1500 °C in air, the SiC-made fixture resulted strongly contaminated by a glassy phase exuded by the tested bar, as depicted in Fig. 7. The exuded glassy product was collected and sampled by FESEM-EDS: the analyses revealed the presence of WO3 with B and C traces. In spite of that, values of great significance were obtained after bending tests at 1500 °C in air: the flexure strength retention of ZS0-15WC, ZS3-5WC, and ZS15-5WC were 94, 113, and 78% of the RT average values, respectively, see Table 5. From Table 5, a tendency for the high-temperature flexure strength behavior can be inferred. Recently, Zou et al. [10] reported that grain sliding and formation of cavities occur preferentially at SiC  Fig. 5. Thermochemical equilibria calculated at 0.001 bar of isobaric pressure representative of the HP vacuum level and initial moles of starting reactants equal to: 1 ZrB2, 1 WC, 0.1 ZrO2 and 0.1 B2O3.  Fig. 6. Reaction products vs. temperature between WC (1 mol) and SiO2 (1 mol) at 0.001 bar of isobaric pressure.  \\x0c\", \"Table 5 Relative density (RD⁎), flexure strength at room temperature (σRT) and at 1500 °C (σ1500) of ZrB2 ceramics initially batched with WC or WSi2.  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  Label  ZS0-15WC  ZS3-5WC ZS15-5WC ZS20-5WC ZS0-15WS ZS15-5WS  pw: present work. only 1 bar tested.  a  Initial secondary phases (vol%)  15 WC  3 SiC, 5 WC 15 SiC, 5 WC 20 SiC, 5 WC 15 WSi2 15 SiC, 5 WSi2  RD⁎ (%)  96.5  N 99.9 93.8 N 99 N 99 N 99  Method  4-pt., air 4-pt., Ar 4-pt.,air 4-pt.,air 3-pt., Ar 4-pt., air 4-pt., air  σRT (MPa)  631 ± 106  -  628 ± 76 543 ± 41 ~ 600 641 ± 19 754 ± 104  σ1500°C (MPa)  596 ± 53 715a 712 ± 23 426 ± 15 ~ 650 537 ± 16 308 ± 43  403  Ref.  pw pw pw pw [10] [19] [33]  locations leading to a degradation of the high-temperature mechanical performances. Surprisingly, ZS15-5WC, with an overall residual porosity of 6 vol%, resulted in higher strength than a fully dense analogous ceramic initially batched with 5 vol% WSi2 instead of WC (see composition labeled ZS15-5WS in Table 5). On the other side, the lack of SiC in a fully dense ZrB2-based matrix enabled to retain the RT values also at 1500 °C, although the traumatized appearance of the bar (Fig. 7). In fact, using 15 vol% of solely WSi2 or WC as starting second phases (ZS0-15WS and ZS015WC, respectively), similar effects on the high-temperature flexure strength were recorded, presumably in relation to common features in the sintered microstructures, i.e. the (W,Zr)B phase and a Wenriched rim around ZrB2 grains. Worth of mention is also the ZS015WC ceramic tested at 1500 °C in a partially protective Argon environment, unfortunately just one bar was available for this preliminary test, but its flexure strength achieved 715 MPa. Average flexure strength around 710 MPa measured at 1500 °C for the ZS3-5WC ceramic was also of notable significance. To the authors' best knowledge, values above 700 MPa were never measured at these temperatures in air. Interesting values were reported for ZS20-5WC (see Table 5), although in Argon and using a 3-pt bending configuration [10]. Another confirmation of the potential of the ZrB2-WC-SiC system to withstand external thermal loads is given by the linearity of the load vs displacement curves at 1500 °C during flexure test (not shown), even for ZS15-5WC in presence of a significant amount of SiC and 6 vol% porosity. The preliminary authors' favorite hypothesis to explain the impressive hightemperature response of the ZrB2-WC-SiC relies on the role of the Wenriched rims. Previous investigations by transmission electron microscopy clearly showed that dislocation networks were generated in ZrB2 grains after bending test at high temperature to progressively accommodate the stress induced by the load [34]. In analogous cases, it was assessed that dislocation chains at the core/rim interfaces develop already after sintering upon cooling, owing to the core/rim thermal expansion mismatches [19,35]. The partial absorption of mechanical energy during the flexural test by dislocation flow and multiplication could contribute to the strength retention at high temperature of ZrB2 ceramics doped with W. An increased dislocations density sets favorable conditions for the creation of dislocation networks, i.e. the creation  of sub-grains, resulting in an overall decrease of the flaws size, from the micrometric scale of the original grains, to the nanometric scale of the dislocations tangles. In addition, the achievement of rim-to-rim grain boundaries depleted of foreign oxide compounds, thanks to reactions catalyzed by the addition of WC, also contributes to maintain outstanding flexure strength at high temperature. Dedicate studies on the high temperature strength behavior of this system are currently in progress.  3.5. Oxidation behavior  3.5.1. Mass change, thermodynamics, and microstructure modification upon oxidation  The exposure of the sintered ZrB2 ceramics, with a native grey-dark visual appearance, to an oxidizing environment for 15 min at 1500 °C in static air induced a color change to dark grey, indication of glass formation, or to whitish-yellow, signature of formation of mainly zirconium oxide. The oxidized samples, displayed with representative insets in Fig. 8, look like darker with isolated bubbles when SiC is present, Fig. 8c,e, while the resulting aspect of the SiC-free ZS0-15WC specimen is characterized by a yellowish and rough oxide, Fig. 8a. Only in this last case, the external scale detached from the inner un-oxidized substrate and spalled off upon rapid cooling (inset of Fig. 8a). The specific mass changes are reported in Table 6: the lowest variation, 3.40 ± 0.05 mg/cm2, was registered for the material containing 15 vol% SiC (i.e. ZS15-5WC) while increasingly larger mass variations were measured for ZS3-5WC (5.35 ± 0.05 mg/cm2) and the SiC-free ZS0-15WC. For this third ceramic, due to partial detachment of the external oxide scale, the measured mass change, considered not reliable, is not reported in Table 6. In all cases, the prevailing crystalline phase identified by XRD on the external oxidized surface was indexed as monoclinic ZrO2, with tetragonal ZrO2 in traces. Peaks of t-WB, in lower intensities compared to the as-sintered microstructure, were still indexed in ZS3-5WC and ZS155WC, i.e. when the silica glass formed on the external surfaces acted as partial inhibitor of the oxygen diffusion into the inner bulk. On the other hand, c-(Zr,W)C and ZrB2 peaks completely disappeared from the XRD patterns of the three oxidized ceramics.  Fig. 7. (a) Flight view and (b) lateral view of the SiC-made fixture after bending test in air at 1500 °C using bars from the ZS0-15WC ceramic.  \\x0c\", '404  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  Fig. 8. FESEM micrographs from representative areas of (a, c, e) the external surface and (b, d, f) fractured cross-section of the ZrB2 ceramics after oxidation at 1500 °C for 15 min. (a-b) ZS015WC, (c-d) ZS3-5WC, (e-f) ZS15-5WC. The visual appearance of the coupons upon air-quenching is shown in the insets (upper right corner a, c, e). Reaction scales are also outlined (dotted lines in b, d, and f).  Based on the expression (2):  I=I0 ¼ exp −μ \\x01 ρ \\x01 τ  ð  Þ  ð2Þ  that describes the attenuation intensity (I/I0) of the initial CuKα X-ray radiation (I0, 30° of glancing angle) at different depth (τ) of a solid specimen with specific density (ρ) and mass attenuation coefficient (μ), scales of 10 μm and 70 μm thick made of only zirconia or silicon oxide, respectively, were calculated being necessary to attenuate 99% of the maximum initial intensity I0. It follows that, the characteristic XRD peaks of t-W boride detected in the case of the ZS3-5WC and ZS155WC oxidized samples, were generated inside the reaction scale (see Fig. 8d,f), not from the un-oxidized bulk.  Table 6 Post-oxidation microstructural features of ZrB2 ceramics after exposure to air at 1500 °C for 15 min: specific mass change Δm/A, (± 0.05 mg/cm2), indexed crystalline products by XRD acquired on the surface and indicative thickness of the external silica-based layer (t1) and of the inner reaction scale (t2).  Δm/A  (mg/cm2)  n.a. 5.34 3.40  Indexed crystalline phases by XRD  Major  m-ZrO2 m-ZrO2 m-ZrO2  Minor  t-ZrO2 WB, t-ZrO2 WB, t-ZrO2  t1  (μm)  -  5 7  t2  (μm)  220-230 85 45  Label  ZS0-15WC ZS3-5WC ZS15-5WC  The reason of the residual retention of WB in the reaction scale was investigated exploring the thermodynamically favorable reactions of an idealized ZrB2-SiC-WB-ZrC system progressively attacked by an increasing supply of oxygen. Thermochemical equilibria computed at 1500 °C and 1 bar of isobaric pressure revealed that WB is the less inclined phase to react with oxygen, compared to ZrC, ZrB2, and SiC, Fig. 9a. Most importantly, at the early stages of the oxidation attack, SiC tends to react with oxygen faster than WB. It follows that WB phase can take advantage from the oxidation of SiC, because the latter phase promptly starts yielding silica glass which protects the inner volume before WB itself begins to oxidize. In the SiC-free ZS0-15WC ceramic, no protection against the inward diffusion of oxygen was afforded by the silica-based glassy coating and the diboride matrix readily transformed into zirconia, which is known to behave only as a moderate barrier to the inward diffusion of oxygen. In order to free our computations from the variables of input amounts and equilibrium pressure, the predominance diagram is a useful tool to better define stability areas of condensed phases vs temperature and/or chemical potential for selected elemental system, W-B-O in the present case. Fig. 9b displays that, at 1500 °C, metallic W can exist when the oxygen partial pressure is very low, like inside the ZrO2 grains, while, with increasing oxygen partial pressure, the formation of WO2 and WO3 is instead favorable. FESEM images of the external surfaces reported in Fig. 8a-e, show that, in absence of SiC, the external oxide scale is wavy and rough,  \\x0c', 'F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  405  Fig. 9. (a) Thermochemical equilibria calculated at 1500 °C, 1 bar of isobaric pressure: starting moles of reactants were 1 ZrB2, 1 WB, 1 ZrC, and 0.5 SiC. (b) Predominance diagram for W-O- B system calculated at 1500 °C.  while, when 3 or 15 vol% SiC is present, the external surface takes a smooth glossy appearance because covered by a continuous silicabased glass. White features consisting of zirconia in form of isolated or agglomerated crystals decorate the external glassy coating, Fig. 8c,e. Representative views of the fractured cross-sections of the oxidized samples are shown in Fig. 8b-f. As first observation, it can be seen that the absence of SiC led the oxidation front to advance deeper into the inner bulk. The final effect was the creation of an irregular oxide layer, whose representative thickness was estimated to span between 220 and 230 μm, see Table 6. Adding SiC, the ZrB2 ceramics are much better  protected against oxidation at the temperature tested, 1500 °C, indicating that a sealing outer glassy scale is more efficient in slowing down the advancement of the oxidation front into the inner bulk, even in presence of 6 vol% of residual porosity, like in the ZS15-5WC ceramic. Indeed, the representative thickness of the modified scale passes from around 80 μm, for 3 vol% of SiC, to 45 μm for 15 vol% of SiC. In detail, in Fig. 10a, the outermost scale of the ZS15-5WC ceramic is based on silica glass where also traces of W were detected by FESEMEDS (Fig. 10, spectrum 1). Encased bright contrasting spots are also visible inside the glassy layer which, in turn, is crossed by zirconia pillars  Fig. 10. Details from the fractured cross-section of ZS15-5WC after oxidation at 1500 °C showing (a) the morphology of the outermost silica-based layer, (b) the consumption of SiC, (c) the presence of frozen W nano-beads onto zirconia grains, and d) the coherent growth of zirconia at the diboride/oxide interface. The lower row shows FESEM-EDS spectra acquired on the numbered spots.  \\x0c', \"406  F. Monteverde, L. Silvestroni / Materials and Design 109 (2016) 396-407  outcropping the external surface. Underneath, rounded zirconia grains with bright discrete pockets of the native partially oxidized (W,Zr)B can be found, Fig. 10a. Right in this region SiC is no longer present, while free C and SiO2 are formed at its place, Fig. 10b, spectrum 2. Moving inwards and deeper into the bulk, the original SiC particles progressively can be found again with un-oxidized (W,Zr)B. Throughout the modified scale, moving from the oxide/boride interface (Fig. 10d) up to the surface (Fig. 10a) rounded nano-sized particles (hereafter indicated as nano-beads) of elemental W, or W oxides, with a bright appearance were identified and analyzed (Fig. 10b,c, spectra 3  and 4). Similar “W-O (and possibly Hf) spherical grains” were observed  by Carney and co-authors [13] on HfB2-SiC-WC materials oxidized for 30 min at 1600 °C in static air. The yield of elemental W, in the present case frozen as nano-bead upon air quenching, is in agreement with the thermodynamic predictions: based on the outputs in Fig. 9, the oxidation of WB appears one reaction that provides elemental W. Actually, the calculations did not include the diboride rim, i.e. (Zr1− xWx)B2, as initial entry, and its W entrapped (see Fig. 2) because not present in the available phase database. In essence, the rim is the largest widespread reservoir for W in the hot-pressed microstructure. According to reactions (3-5), W present in the rims is first released as element and subsequently converted into solid/liquid W oxides.  ð  Zr1−xWx  ÞB2 þ 2:5−x  ð  Þ O2 gð Þ→ 1−x Þ ZrO2 þ x W sð  ð  Þ þ B2O3  ð  Þ ð3Þ  l; g  W sð  Þ þ O2 gð  Þ→WO2  sð  Þ  WO2  sð  Þ þ 0:5 O2 gð Þ→WO3  lð Þ  ð4Þ  ð5Þ  At 1500 °C, liquid WO3 alone, i.e. not embedded into boro-silicate glass, is not stable anymore, but tends to volatilize forming complex gaseous W oxides, like W3O9, W2O6, W4O12, and W3O8 [36]. On one hand, the lack of compactness followed by the oxide scale spall off (see Fig. 8) finds its original cause in the extensive release of volatile W oxides. On the other hand, when WO3 is readily trapped within the (boro)-silicate glass, the disruptive consequences of the evolving W oxide are minimized. The reason why a dense population of W nanobeads survived upon oxidation and fast air quenching is very likely related to their position, well nestled inside ZrO2 grains where the oxygen partial pressure does not exceed 10− 8 Pa, see Fig. 9b. Lastly, the interface between the reaction front and the un-oxidized bulk is regular, well adherent, and without cracks, Fig. 10d.  3.5.2. Oxidation mechanism  The main mechanism responsible of the improved oxidation resistance in the ZrB2-SiC-WC systems compared to the SiC-deficient one can be purely related to the ability of the in situ formed external silica-based coating to limit oxygen inward diffusion. (Zr,W)B2 rims, as well as (Zr,W)C and (W,Zr)B phases, are present in the sintered bulk and their resulting oxidation products consist of ZrO2(s), W(s), WO3(l), and amorphous B2O3(l,g). It follows that the glassy silica can change its own composition embedding foreign elements such as Zr, W, and B into its glass network. If the availability of W in the silica glass, once modified its actual viscosity, is capable of slowing down the effective diffusion rate of oxygen is still matter of current investigations [33]. Another beneficial mechanism that could take part to the bulk protection against oxidation concerns the modification of the characteristics of the forming zirconium oxide. In the conditions of a reduced oxygen partial pressure, zirconium oxide is known to form and evolve featured by vacancies [37], in a more pronounced density when the reaction front is deep close to the diboride/oxide interface. In this respect, it should be underlined that the base matrix of all these ceramics is characterized by a homogeneously distributed mixed diboride, (Zr,W)B2, therefore an intimate coexistence of W in the zirconia crystal lattice, at least in the early oxidation stage, very likely took place. The solubility  of W in ZrO2 is very limited, below 2 at.% [38], therefore the threshold limit is quickly achieved resulting in the expulsion of W that tends to aggregate as nano-beads across the whole zirconia scale, Fig. 10c. However, the initial presence of guest W cations with higher valence than Zr in the zirconia hosting lattice can act as vacancy filler for the ongoing formation of zirconia deep inside the reaction scale. Another feature worthy of mention is the formation of a coherent and sharp oxidation front, which has a regular thickness and no cracks are present at the oxide/boride interface, Fig. 8b-f. While for other ZrB2-based materials, containing Si3N4 [39] or other transition metal silicides (Zr, Mo, Ta) [21], the oxidation front proceeded along grain interfaces even in absence of amorphous films, in the present case the diboride grains were entirely attacked trans-granularly, as depicted in Fig. 10d. This behavior was related once again to the better oxidation resistance of W-containing borides, constituting the adjacent shells of the grains, over pure ZrB2, constituting the cores.  4. Conclusions  ZrB2 ceramics containing varying amount of SiC (0, 3, 15 vol%) and WC (5, 15 vol%) were processed by hot-pressing and relative density above 94% were reached. The resulting microstructures were characterized by highly refractory phases, solid solutions and grain junctions depleted of amorphous glassy phases: this enabled to improve the state of art of ZrB2 ceramics in terms of high-temperature capabilities. The simultaneous addition of SiC (3 vol%) and WC (5 vol%) led the hot-pressed ZrB2 ceramic to exhibit room temperature average strength of 630 MPa that increased up to 720 MPa upon testing at 1500 °C in air. The impressive mechanical response of this system in an oxidizing environment was explained in terms of a combination of favorable conditions:  i)  ii) iii)  iv)  the extensive formation of core-rim substructures, the latter constituted by (Zr,W)B2 solid solutions; rim-to-rim grain boundaries depleted of secondary phases; the formation of continuous glassy layer upon oxidation at 1500 °C, that effectively healed surface flaws; the presence of WB phase, more oxidation resistant than ZrB2, which slowed down oxygen diffusion through grains interfaces.  Finally, the accumulation of dislocations at the core-rim interfaces was indicated as a favorable condition for plastic deformation to accommodate the stress under loading at high temperature which leads the material to retain its thermo-mechanical performances.  Acknowledgments  The research leading to these results has received funding from the European Community's Seventh Framework Programme (FP7/20112014) under grant agreement LIGHT-TPS No. 607182. The authors are grateful to D. Dalle Fabbriche and C. Melandri (ISTEC-CNR) for their support in performing hot-pressing and mechanical testing, respectively.  References  [1]  E. Wuchina, E. Opila, M. Opeka, W. Fahrenholtz, I. Talmy, UHTCs: ultra-high temperature ceramic materials for extreme environment applications, Electrochem. Soc. Interface (Winter 2007) 30-36. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [3] D. Sciti, L. Silvestroni, V. Medri, F. Monteverde, Sintering and densification mechanisms of ultra-high temperature ceramics, in: W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. 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  "_id": 30,
  "PDF": "Comparative study on microstructure and oxidation behaviour of ZrB2-20 vol_ SiC ceramics reinforced with Si3N4-Ta additives.pdf",
  "Text": "['Journal of Alloys and Compounds 797 (2019) 92e100  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : h t t p : / / w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  Comparative study on microstructure and oxidation behaviour of ZrB2-20 vol% SiC ceramics reinforced with Si3N4/Ta additives  Brahma Raju Golla*, Sravan Kumar Thimmappa  Metallurgical and Materials Engineering Department, National  Institute of Technology, Warangal, 506 004,  India  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 28 March 2019 Received in revised form 3 May 2019 Accepted 8 May 2019 Available online 9 May 2019  Keywords:  ZrB2 Si3N4 Ta Spark plasma sintering Microstructure Oxidation  Multi-stage Spark Plasma Sintering (MS-SPS) of ZrB2-20 vol% SiC-10 vol% Si3N4 (ZSS) and ZrB2-20 vol% SiC-10 wt% Ta (ZST) composites were carried out at 1900 \\x0e C for 3 min to densify the ZrB2-based composites. The microstructure of ZSS composite consisted of new secondary phases (ZrO2, BN, ZrN) along with SiC and ZrB2. Interestingly, ZST sample exhibited core-shell/rim structure and it comprised of ZrB2 core, (Zr,Ta)B2 rim, SiC, ZrO2 and (Zr,Ta)C phases. The presence of new phases and annihilation of Si3N4 or Ta additive in both the sintered samples was clear indication of involvement of sintering reactions. The SEM of cross sectional samples revealed presence of three distinctive layers for the ZrB2 samples after the isothermal oxidation at 1600 \\x0e C for 10 h. In particular, no SiC depleted layer was observed for ZSS and its presence was evident in ZST composite. The weight gain (varied between 15.25 and 16.66 mg/cm2) of the ZrB2 composites was comparable and significant difference in the oxide layer thickness (changed from 255 to 476 mm) was noticed. Overall the ZST sample offers better oxidation resistance in view of protective nature of its oxide layer.  © 2019 Elsevier B.V. All rights reserved.  1.  Introduction  In the last two decades, Ultra-High Temperature Ceramics (UHTCs) have drawn considerable research interest due to the increasing demand for hypersonic re-entry vehicles, thermal protection systems (TPS) and energy applications etc. [1e13]. Among the UHTCs, Zirconium diboride (ZrB2) is the most widely investigated material as it exhibits good combination of properties: low density, high hardness, strength and elastic modulus, excellent thermal and electrical conductivity and very good high temperature properties [2,3,14]. However, the applications of monolithic ZrB2 is restricted due to its poor densification/sinterability, moderate fracture toughness and low oxidation resistance at above 1100 \\x0eC [1e3,6,14e19]. The use of sintering additives (both metallic and non-metallic) and various sintering techniques (conventional and advanced) have been employed to improve the densification and properties of ZrB2 [2,3,14,15]. Non-metallic additives such as SiC, C, B4C, WC, ZrC, AlN, Si3N4, MoSi2, TaSi2 and ZrSi2 etc. have been used to overcome the limitations of ZrB2 [2,3,6,14,15,17e20]. It was well reported that the  * Corresponding author. E-mail addresses: gbraju@nitw.ac.in, gbraju121@gmail.com (B.R. Golla).  https://doi.org/10.1016/j.jallcom.2019.05.097 0925-8388/© 2019 Elsevier B.V. All rights reserved.  sinter-additives enhance the densification of borides primarily by removing the surface oxide impurities of starting boride powders. Among the non-metallic additives, SiC has been extensively selected for ZrB2, since the incorporation of SiC enhances the sinterability, inhibits grain growth and improves mechanical, oxidation and other high temperature properties of ZrB2. The addition of 20 vol% SiC to ZrB2 was proved to be optimal composition for hypersonic applications as it exhibited good combination of properties [2,21,22]. In the present work, ZrB2-20 vol% SiC (ZS) was selected as a base line material to understand the effect of Si3N4/Ta additives on its microstructure and properties. In fact, such studies were not explored much in the literature. Metallic additives (generally up to 10 wt%) such as Ni, Cu, Co, Cr and Fe etc. were added to avoid processing difficulties, reduce the sintering temperature and improve (room temperature) mechanical properties of ZrB2 [2,3]. Hence, the ZrB2 with metallic additives cannot be used for high temperature applications. However, in recent times, the effect of transition metals (Hf, Nb, W, Ti, Y etc.) on densification, microstructure and thermal properties of ZrB2 has been explored [23e25]. It has been reported that the addition of transition metals to the diborides modifies its microstructure in the form of core-shell/rim structure and improve high temperature mechanical properties. In the present work, multi-stage Spark Plasma Sintering (MS \\x0c', 'B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  93  1900 \\x0eC for a short duration of 3 min under the uniaxial pressure of 50 MPa. Isothermal oxidation tests were performed at 1600 \\x0e C for a duration of 10 h in air to assess thermal stability of the ZrB2 composites. The oxidation resistance of samples was evaluated by measuring weight and oxide layer thickness. The microstructural studies (surface and cross-sectional microstructure analysis) were performed using XRD and SEM-EDS.  2.  Experimental procedure  Both the ZrB2-20 vol% SiC-10 vol% Si3N4 (ZSS) and ZrB2-20 vol% SiC-10 wt% Ta (ZST) composites were prepared using commercially available ZrB2 (purity >98.2%, H.C. Starck Grade B GmbH and Co., Goslar Germany), SiC (purity >99.8%, Alfa Aesar), Si3N4 (purity >99.8%, Alfa Aesar) and Ta (purity >99.99%, India) powders. The powders in appropriate amounts were weighed and mixed in a planetary ball mill (Fritsch Pulverisette 6, Germany) using silicon nitride (Si3N4) milling media in toluene for 6 h at 200 rpm with balls to powder weight ratio at 2:1. After mixing the powder compositions, the slurries were dried in a rotary evaporator (at ~98 \\x0e C) to remove toluene and minimize the particles agglomeration. The sintering of powders was carried out at different stages (multi-stage) using SPS (Model: MS-SPS 25-10, GT Advanced Technologies, USA) at a temperature of 1900 \\x0eC for 3 min under \\x003 Pa) using graphite die and 50 MPa pressure in a vacuum (10 punches. More details about the MS-SPS schedule can be found elsewhere [30]. After the MS-SPS, the specimens (15 mm in diameter and thickness of 3e4 mm) were subjected to grinding and  Fig. 1. XRD patterns of ZrB2, SiC, Si3N4 and Ta starting powders. All consisted of respective phase without any secondary phases.  the powders  SPS) was used to densify ZS ceramics reinforced with Si3N4/Ta. In fact, limited or no research is available on understanding microstructure and oxidation resistance of ZrB2-SiC ceramics reinforced with Si3N4/Ta [23e29]. A comparative study was made with respect to microstructure and oxidation behaviour as the addition of Si3N4 and Ta additives to ZrB2-20 vol% SiC resulted absolutely different microstructure. The MS-SPS experiments were carried out at  Fig. 2. SEM micrographs of (a) ZrB2, (b) SiC, (c) Si3N4 and (d) Ta starting powders.  \\x0c', \"94  B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  Table 1 Densification, microstructure and oxidation results of ZrB2-SiCSi3N4/Ta composites after MS-SPS at 1900 \\x0e C, 50 MPa, 3 min. The oxidation tests were carried out at 1600 \\x0e C for 10 h in air. (RDRelative density).  Sample Composition  Sample designation  Exp. Density (g/cc)  RD (%)  Microstructural Phases after sintering  Weight gain (mg/ cm2)  Oxide layer thickness (mm)  Microstructural Phases after oxidation  ZrB2-20 vol% SiC-10 vol% Si3N4 ZrB2-20 vol% SiC-10 wt% Ta  ZSS  ZST  5.11  5.87  98.7  ZrB2, SiC, ZrN, BN, ZrO2  99.8  ZrB2, (Zr,Ta)B2, SiC, ZrO2, (Zr,Ta)C  15.25  16.66  476 ± 5  255 ± 4  SiO2, ZrO2  SiO2, ZrO2, TaZr2.75O8  element) by heating samples (2 \\x02 4 \\x02 8 mm3) at 1600 \\x0e C for 10 h in stagnant air. The samples were cleaned in acetone thoroughly and were dried. Careful weight measurements of the samples were performed before and after the oxidation test. The cleaned coupons were placed in a pure alumina crucible with minimal contact area and then the crucible was placed at the center portion of furnace. The microstructure of ZrB2 starting materials, sintered samples and oxidation tested samples was characterized using scanning electron microscope (SEM, TESCAN VEGA 3 LMU) coupled to an energy dispersive spectroscope (EDS, Oxford Instruments). The crystalline phases of the powders and sintered samples were characterized by X-ray diffraction (XRD, PANalytical X'PertPro, Holland, CuKa ¼ 1.5405 Å). Both surface and cross-section microstructural characterization of the oxidized samples were performed using SEM-EDS and XRD.  3. Results and discussion  3.1.  Starting materials  The XRD of starting powders (ZrB2, SiC, Si3N4 and Ta) is shown in Fig. 1. It consists of respective crystalline phases of the powders and no other additional phases were observed within the detection limit of XRD. Fig. 2 depicts the microstructure of all the starting powders that were used to produce ZSS and ZST composites. The ZrB2 powders are of platelet shape with its particle size varying in the range 0.5e4.0 mm (Fig. 2a). Whereas the SiC powders are fine and its particle size observed to vary from 0.3 to 2.6 mm (Fig. 2b). The Si3N4 powders are of much finer, irregular and its particle size vary in the range 0.1e1.6 mm. The Ta powders are of mixed shape and its particle size vary between 0.6 and 7.0 mm.  3.2. Microstructure of sintered ZSS samples  The experimental density of ZSS composite was measured to be 5.11 g/cc, and its relative density (RD) was about 98.73% (see Table 1). The XRD of ZSS reveals the presence of ZrN, BN and ZrO2 along with the ZrB2 and SiC major constituent phases (Fig. 3a). Remarkably, no traces of Si3N4 could be observed in the XRD of ZSS. The presence of new phases (ZrN, BN and ZrO2) and absence of Si3N4 clearly indicate involvement of sintering reactions. The possible sintering reactions are as the following.  2B2O3 þ Si3N4 ¼ 4BN þ 3SiO2  Si3N4 ¼ 3Si þ 2N2  2ZrB2 þ 3Si þ 3N2 þ 3C ¼ 2ZrN þ4BN þ 3SiC  (1)  (2)  (3)  The reasons for the involvement of sintering reactions and its thermodynamic feasibility were reported in one of our recent works [30]. The presence of such new phases was also proved with  Fig. 3. XRD patterns of (a) ZSS and (b) ZST samples after MS-SPS at 1900 \\x0e C, 50 MPa, 3 min.  polishing and cleaned with acetone using ultrasonic bath. The experimental densities of samples were measured by the Archimedes method. The relative density (RD) or densification of samples was estimated by determining the residual porosity present in the SEM micrographs using ImageJ software package (ImageJ 1.51j8, National Institute of Health, USA). The isothermal oxidation tests of the ZrB2 composites was carried out using box type furnace (Nabertherm, MoSi2 heating  \\x0c\", 'B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  95  Fig. 4. SEM microstructure of MS-SPSed ZSS composite with detailed EDS analysis, showing the various phases: ZrB2, SiC, ZrN and ZrO2.  Zr-N-B ternary phase diagram. The SEM image indicates presence of different contrasting phases in the MS-SPSed ZSS (Fig. 4). The SEM-EDS confirms bright contrasting grains as ZrB2, the grey phase as SiC and the dark phase as ZrN. The irregular bright contrasting phase (small size) at the grain boundaries was identified as ZrO2. All the secondary phases were well dispersed in the ZrB2 matrix. The grain size of sintered ZrB2 is very much comparable with the initial particle size of ZrB2 powders. In the literature, formation of new phases after sintering of ZrB2Si3N4 ceramics has also been reported by several researchers. They also attributed the formation of new phases to reaction between Si3N4 and oxides of boron and the silicon oxides present on Si3N4 powders and ZrB2. For example, the formation of BN, ZrSi2, ZrN and ZrB2 phases were reported in hot pressed ZrB2-(5-35)vol% Si3N4 composites [28]. Monteverde et al. [29] reported the presence of grain boundary phases (BN, ZrO2, ZrSi2 and B-N-O-Zr-Si glassy phase) for ZrB2-2.5 wt% Si3N4 composites. In another work, presence of BN, ZrO2 and glassy phase (B-N-O-Zr-Si) were observed with hot pressed ZrB2-5 vol% Si3N4 ceramics [27].  Si3N4. Fig. 3b presents the XRD of ZST sample after sintering. The presence of ZrB2, (Zr,Ta)B2, SiC major phases along with ZrO2 and (Zr,Ta)C minor phases is evident from the XRD. These observations indicate that both ZSS and ZST consists of new crystalline phases after MS-SPS. Fig. 5 reveals core and rim structure of ZST and is very different from ZSS. The SEM-EDS analysis confirms the microstructure consists of ZrB2 (core), (Zr,Ta)B2 (rim), SiC and bright (Zr,Ta)C solid solution phase at grain boundary. It is expected that tantalum separates the ZrB2 matrix and the Ta dissolves into ZrB2 matrix and forms the solid solution rim phase. The core-rim structure was observed for ZrB2 reinforced with different silicides, due to the reaction between the silicide and surface oxides of ZrB2 [17,31,32]. For instance, the ZrB2 core and (Zr,Ta)B2 rim structure was reported for ZrB2-15 vol% TaSi2 and ZrB2-20SiC-(010)TaSi2 (vol%) composites [31,32]. On the other hand, only the presence of (Zr,TM)B2 solid solution phase was noticed when ZrB2 was added with small amount (1-1.2 wt%) of transition metals (TM) [23]. It is mainly because of dissolution of the transition metal in to ZrB2.  3.3. Microstructure of sintered ZST samples  3.4. Oxidation of ZSS and ZST composites at 1600 \\x0e C  The density of ZST composite was observed to be slightly higher (5.87 g/cc), and it was almost fully densified (99.63%) (see Table 1). It appears that Ta is aiding in the densification of ZS than that of  Table 1 records the weight gain and oxide layer thickness of the ZrB2 samples after oxidation at a temperature of 1600 \\x0e C for 10 h. It can be noticed that the weight gain of samples varied narrowly  \\x0c', '96  B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  Fig. 5. SEM microstructure of MS-SPSed ZST composite with detailed EDS analysis, showing the phases: ZrB2, SiC, (Zr,Ta)B2 and (Zr,Ta)C solid solution phase.  between 15.25 and 16.66 mg/cm2 and the oxide layer thickness changed considerably from 255 to 476 mm. In particular, the weight gain is slightly high and the oxide layer thickness significantly low for ZST in comparison to ZSN. The XRD of ZrB2 sample surfaces after oxidation at 1600 \\x0e C for 10 h is shown in Fig. 6. In case of ZSS, the presence of crystalline ZrO2 and amorphous SiO2 (peak broadening at about 22\\x0e ) phases was noticeable. However, the presence of ZrO2 and TaZr2.75O8 crystalline phases along with amorphous SiO2 was observed in ZST. The SEM-EDS of oxidized surfaces of the ZrB2 samples is presented in Fig. 7. The microstructure of ZSS reveals the manifestation of irregular shaped (dendritic and spherical) ZrO2 in SiO2 (dark phase) matrix. The size of ZrO2 ranges from submicron to 6 mm. The EDS analysis also confirms the presence of ZrO2 and SiO2 phases on the oxidized surface of ZSS sample (Fig. 7a). The microstructure of ZST consists of ZrO2, TaZr2.75O8 and SiO2 phases (see Fig. 7b). These phases were confirmed by SEM-EDS, in particular, the morphology of ZrO2 appeared to be spherical (fine and coarse). A careful analysis of coarse ZrO2 phase reveals the coexistence of bright contrasting TaZr2.75O8 rim phase around grey contrast ZrO2 core. It appears that TaZr2.75O8 grows on ZrO2 through the diffusion of Ta. The dark contrasting matrix phase was identified as SiO2. The cross-sectional SEM-EDS analysis of ZSS and ZST samples after oxidation is presented in Fig. 8. From Fig. 8, it is obvious that both samples consist of three distinctive layers after oxidation. In case of ZSS, the oxide layer consisted of dense outer thick layer of SiO2, intermediate layer of ZrO2-SiO2 and unreacted bulk (Fig. 8a).  The ZST sample composed of outer dense SiO2 layer, SiC depleted layer and unreacted bulk. All these oxide layers presence was confirmed with EDS analysis. In addition to the striking differences in the oxide layer composition, considerable difference in the thickness of layers was also clearly observed for both the ZrB2 samples. These microstructural differences indicate the addition of Si3N4 and Ta controls the oxidation of ZS via different mechanisms.  3.5. Oxidation mechanism in ZSS  The oxidation of ZrB2 in air is more prominent at high temperatures (more than 1200 \\x0eC) and it leads to the formation of ZrO2 and B2O3 phases on its surface. However, B2O3 evaporates in the temrange of 1100e1400 \\x0e C and formation of porous ZrO2 perature lowers the oxidation resistance of ZrB2 [18,27]. However, the oxidation resistance of ZrB2 can be improved with the addition of SiC/Si3N4 as it forms protective borosilicate glass or SiO2 rich oxide layer on the surface. The oxidation of ZSS sample is expected to take place through the following oxidation reactions.  ZrB2(s)þ5/2O2(g) ¼ ZrO2(s)þB2O3(l)  B2O3(l) ¼ B2O3(g)  2ZrN(s)þ2O2(g) ¼ 2ZrO2(c)þN2(g)  (4)  (5)  (6)  \\x0c', 'B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  97  cited that the SiC depleted layer may reduce the oxidation resistance of ZrB2.  3.6. Oxidation mechanism in ZST  In case of oxidation of ZST, as SiC depleted layer was observed it may be possible that active SiC oxidation reactions might have taken place along with the other reactions. Hence reactions (4), (5), (8), (9), (10) and the following reactions involve in the oxidation of ZST.  (Ta,Zr)B2(s) þ O2(g) / TaZr2.75O8(s)  (Ta, Zr)C(s) þ O2(g) / TaZr2.75O8(s)  (11)  (12)  Although the SiC depleted layer was observed in ZST sample, there is no indication of any defects at the SiO2 layer (see Fig. 8). Several researchers have observed the presence of SiC depleted layer in ZrB2-SiC composites [18]. Also SiC depleted layer was reported for TaSi2 and TiB2 reinforced ZrB2-SiC composites [18,19,31,33e36]. As mentioned above, the SiC depleted layer lowers the oxidation resistance of ZrB2 composites. However, in the present work, despite the presence of SiC depleted layer in ZST, the sample exhibited comparable weight gain and considerable low oxide layer thickness than ZSS. Even the presence of core-rim structure of boride composites reportedly exhibited good elevated temperature strength properties [37,38]. In view of this, it can be inferred that ZST can be better choice for high temperature applications. Nevertheless, future studies need to be focused on evaluating strength retention capacity of oxidized ZSS and ZST samples.  4.  Conclusions  a ZSS and ZST composites resulted highly dense (more than 98%) ZrB2 materials after MS-SPS at 1900 \\x0e C for 3 min. b A distinctive difference in the microstructure of both the ZSS and ZST composites was observed. In particular, the microstructure of ZSS composites consisted of new secondary phases (ZrO2, BN, ZrN) along with SiC and ZrB2 major phases. On the other hand, ZST composite exhibited core-shell/rim structure and it consisted of ZrB2 core, (Zr,Ta)B2 rim, SiC, ZrO2 and (Zr,Ta)C solid solution phase. c The weight gain of ZSS and ZST composites was comparable and it varied narrowly between 15.25 and 16.66 mg/cm2 and the oxide layer thickness changed considerably from 255 to 476 mm after the isothermal oxidation at 1600 \\x0e C for 10 h in air. d In case of ZSS, the oxide layer consists of dense outer thick layer of SiO2, ZrO2-SiO2 intermediate layer and unreacted bulk. The microstructure of ZST composite composed of outer dense SiO2 layer, SiC depleted layer and unreacted bulk. Overall, the ZST sample offers better oxidation resistance in view of protective nature of its oxide layer. e The presence of core-rim structure of ZST composites is expected to be beneficial for elevated temperature applications. Hence future studies need to be focused on evaluating strength retention capacity of oxidized samples to assess its potential.  Data availability statement  Due to the sensitive nature of the present work, survey respondents were assured raw data would remain confidential and would not be shared.  Fig. 6. XRD patterns of (a) ZSS and (b) ZST samples after oxidation at 1600 \\x0e C for 10 h.  4BN(s)þ3O2(g) ¼ 2B2O3(l)þ2N2(g)  SiC(s)þ3/2O2(g) ¼ SiO2(l)þCO(g)  SiC(s) þ O2(g) / SiO(g) þ CO(g)  SiC(s)þ 2SiO2(l) / 3SiO(g) þ CO(g)  (7)  (8)  (9)  (10)  In the case of ZSS, the oxidation of ZrB2 and ZrN phases causes porous ZrO2. During oxidation of SiC, liquid SiO2 forms and it may fill the pores of ZrO2. The SiO2 layer that forms on the surface of ZrB2 is protective and it inhibits the inward diffusion of oxygen and thus enhances oxidation resistance of ZrB2 ceramics. It was reported that rich SiO2 oxide glassy layer that was formed on pure SiC causes protection up to 1600 \\x0e C [19]. Reactions (9) & (10) represents active oxidation of SiC and it may lead to SiC depleted layer formation. Since no SiC depleted was observed in ZSS, it is possible that active oxidation reactions of SiC (reactions (9) & (10)) might not have taken place. As far as the oxidation of ZrB2-SiC composites is concerned, SiC depleted layer along with SiO2 and unreacted bulk has been reported in most of the earlier studies [6,18,19]. It was also  \\x0c', '98  B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  Fig. 7. Surface morphology of (a) ZSS (b) ZST samples after oxidation at 1600 \\x0e C for 10 h with detailed EDS of various microstructural phases.  \\x0c', 'B.R. Golla, S.K. Thimmappa /  Journal of Alloys and Compounds 797 (2019) 92e100  99  Fig. 8. Cross-sectional SEM of (a) ZSS and (b) ZST samples after oxidation at 1600 \\x0e C for 10 h, showing different stack of  layers with EDS.  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Liu, W. Tian, M. Zhang, X. Jian, L. Yin, L. Zhang, J. Xie, L. Deng, Pursuing enhanced oxidation resistance of ZrB2 ceramics by SiC and WC co-doping, J. Eur. Ceram. Soc. 38 (16) (2018) 5311e5318. [23] D.L. McClane, W.G. Fahrenholtz, G.E. Hilmas, Thermal properties of (Zr,TM)B2 solid solutions with TM ¼ Hf, Nb, W, Ti, and y, J. Am. Ceram. Soc. 97 (5) (2014) 1552e1558. [24] M.K. Dehdashti, W.G. Fahrenholtz, G.E. Hilmas, Oxidation diboride with niobium additions, J. Eur. Ceram. Soc. 33 1591e1598.  of zirconium (10) (2013)  [20]  [21]  [31]  [30]  [29]  [28]  [25] X. Sun, W. Han, P. Hu, Z. Wang, X. Zhang, Microstructure and mechanical properties of ZrB2-Nb composite, Int. J. Refract. Metals Hard Mater. 28 (3) (2010) 472e474. [26] W.M. Guo, L.X. Wu, S.X. Gu, Y. You, H.T. Lin, S.H. Wu, G.J. Zhang, Chemical reactivity of hot-pressed Si3N4-ZrB2 ceramics at 1500-1700 \\x0e C, J. Eur. Ceram. Soc. 35 (11) (2015) 2973e2979. [27] A. Bellosi, F. 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Mater. 160 (2019) 1e4. [38] H.B. Ma, J. Zou, J.T. Zhu, L.F. Liu, G.J. Zhang, Segregation of tungsten atoms at ZrB2 grain boundaries in strong ZrB2-SiC-WC ceramics, Scr. Mater. 157 (2018) 76e80.  [33]  [35]  [37]  [34]  \\x0c']"
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  "PDF": "Comparison of the Oxidation Behaviour of Two Dense Hot Isostatically Pressed Tantalum Carbide (TaC and Ta,C) Materials.pdf",
  "Text": "['J o u r n a l   o f   t h e   E u r o p e a n   C e r a m i c   S o c i e t y   1 7   ( 1 9 9 7 )   1 3 2 5 1 3 3 4   0   1 9 9 7   E l s e v i e r   S c i e n c e   L i m i t e d   P r i n t e d   i n   G r e a t   B r i t a i n .   A l l   r i g h t s   r e s e r v e d   P I I :   S O 9 5 5 2 2 1 9 ( 9 6 ) 0 0 2 3 5 X   0 9 5 5 2 2 1 9 / 9 7 / S 1 7 . 0 0   C o m p a r i s o n   o f   t h e   O x i d a t i o n   B e h a v i o u r   o f   T w o   D e n s e   H o t   I s o s t a t i c a l l y   P r e s s e d   T a n t a l u m   C a r b i d e   ( T a C   a n d   T a , C )   M a t e r i a l s   M .   D e s m a i s o n B r u t ,   N .   A l e x a n d r e   a n d   J .   D e s m a i s o n   L a b o r a t o i r e   d e   M a t C r i a u x   C C r a m i q u e s   e t   T r a i t e m e n t s   d e   S u r f a c e   L M C T S ,   U R A   C N R S   3 2 0     U n i v e r s i t k   d e   L i m o g e s ,   1 2 3   A v e n u e   A l b e r t   T h o m a s ,   8 7 0 6 0   L i m o g e s   C e d e x ,   F r a n c e   ( R e c e i v e d   2 3   J u l y   1 9 9 6 ;   r e v i s e d   v e r s i o n   r e c e i v e d   7   N o v e m b e r   1 9 9 6 ;   a c c e p t e d   1 1   N o v e m b e r   1 9 9 6 )   A b s t r a c t   1   I n t r o d u c t i o n   I s o t h e r m a l   o x i d a t i o n   o f   d e n s e   H I P e d   t a n t a l u m   c a r b i d e   m a t e r i a l s   T a C   a n d   T a z C ,   h a s   b e e n   p e r   f o r m e d   i n   f l o w i n g   o x y g e n   b e t w e e n   7 5 0   a n d   8 5 0 ° C .   T h e   b e h a v i o u r   o f   t h e   t w o   c a r b i d e s :   i . e .   T a C   ( N a C l   t y p e   s t r u c t u r e )   a n d   T a 2 C   ( h e x a g o n a l   t y p e ) ,   i s   c h a r a c t e r i z e d   b y   t h e   g r o w t h   o f   a   n o n p r o t e c t i v e   o x i d e   s c a l e   w h i c h ,   o n   s q u a r e   s e c t i o n   s a m p l e s ,   f o r m s   a   m a l t e s e   c r o s s .   X R a y   d t j i i a c t i o n   a n a l y s i s   h a s   o n l y   s h o w n   t h e   f o r m a t i o n   o f   t a n t a l u m   h e m i p e n t o x i d e   p T a , O , .   T h e   o x i d a t i o n   o f   T a C   p r o c e e d s   b y   a n   i n t e r f a c i a l   r e a c t i o n   p r o c e s s .   F o r   T a , C ,   t h e   m e c h   a n i s m   c o u l d   b e   m o r e   c o m p l e x   d u e   t o   t h e   p r e s e n c e   o f   a n   i n t e r m e d i a t e   o x y c a r b i d e   l a y e r   T a C , O ,   w h i c h   h a s   b e e n   d e t e c t e d   a t   t h e   T a J T a 2 0 ,   i n t e r f a c e .   I n d e e d ,   i n   t h i s   c a s e ,   i t   i s   n o t   p o s s i b l e   t o   e x c l u d e   a   d @ s i o n   l i m i t i n g   p r o c e s s   t h r o u g h   t h i s   o x y n i t r i d e   s u b l a y e r   o f   c o n s t a n t   t h i c k n e s s   w i t h   t i m e .   0   1 9 9 7   E l s e v i e r   S c i e n c e   L i m i t e d .   T r a n s i t i o n   m e t a l s   s u c h   a s   t a n t a l u m ,   h a f n i u m ,   z i r c o n i u m ,   n i o b i u m ,   a s   w e l l   a s   t h e i r   n i t r i d e s   a n d   c a r b i d e s ,   a r e   u n s t a b l e   i n   o x i d i z i n g   a t m o s p h e r e s .   O x i d a t i o n   s t u d i e s   h a v e   b e e n   c a r r i e d   o u t   o n   p o w d e r s ,   w i r e s   o r   p l a t e l e t s ,   b u t ,   f o r   d e n s e   s i n t e r e d   m a t e r i a l s ,   t h e r e   i s   s t i l l   a   c o n s i d e r a b l e   l a c k   o f   k i n e t i c   d a t a .   I n   t h e   p a s t ,   a t t e n t i o n   h a s   c o n c e n   t r a t e d   o n   t h e   t r a n s i t i o n   m e t a l   c a r b i d e s l e 3   w i t h   f e w   w o r k s   b e i n g   c o n c e r n e d   w i t h   t h e   s i n t e r i n g   o f   t a n t a l u m   c a r b i d e   T a C . 6 6   T a n t a l u m   m o n o c a r b i d e   T a C   ( N a C l   t y p e   s t r u c   t u r e )   c o n s t i t u t e s   t h e   m a t r i c e s   o f   h a r d   a l l o y s   f o r   t h e   m a c h i n i n g   o f   m a t e r i a l s .   I n   t h e   l a r g e   h o m o   g e n e i t y   r a n g e   f o r   t h e   f e e   p h a s e   o f   t h e   T a C   s y s t e m ,   e v a l u a t i o n   o f   s u c h   p r o p e r t i e s   a s   m i c r o h a r d n e s s ,   e l e c t r i c a l   r e s i s t i v i t y   o r   s u p e r c o n d u c t i v i t y 7 * 8   e x i s t s .   L e   c o m p o r t e m e n t   a   l b x y d a t i o n   d a n s   l b x y g e n e   d e   d e u x   c a r b u r e s   d e   t a n t a l e   d e n s e s   T a C   e t   T a 2 C   a   t t e   t t u d i e   e n t r e   7 5 0   e t   8 5 0 ° C .   T h e   p h a s e   d i a g r a m   o f   t h e   t a n t a l u m c a r b o n   s y s t e m   s h o w s   a n o t h e r   w e l l d e f i n e d   c a r b i d e ,   T a , C .   T h i s   h e m i c a r b i d e   T a 2 C   h a s   a   s t r u c t u r e   b a s e d   o n   a   h e x a g o n a l   c l o s e s t p a c k e d   m e t a l   l a t t i c e   w i t h   t h e   c a r b o n   a t o m s   f i l l i n g   o n e   h a l f   o f   t h e   o c t a h e d r a l   h o l e s . ‘ , ”   D a n s   l e s   d e u x   c a s ,   l a   c o u c h e   d ’ h e m i p e n t o x y d e   d e   t a n   t a l e   p   T a 2 0 5   s   b u v r e   p r o g r e s s i v e m e n   t   a u   n i v e a u   d e s   a & e s ,   s o u s   l ’ e f f e t   d e s   c o n t r a i n t e s   d e   c r o i s s a n c e ,   e n   f o r m a n t   u n e   c r o i x   d e   m a l t e .   I 1   e n   r e s u l t e   q u e   l ’ o x y d e   n ’ e s t   p a s   p r o t e c t e u r .   C o n c e r n a n t   T a C ,   l e   m t c a n i s m e   p r o p o s e   f a i t   i n t e r v e n i r   u n   r e g i m e   d e   r e a c t i o n   a   l ’ i n t e r f a c e   c a r b u r e o x y d e .   D a n s   l e   c a s   d e   T a z C ,   l e   m e c a n i s m e   e s t   c e r t a i n e m e n t   p l u s   c o m   p l e x e ,   c e c i   e n   r e l a t i o n   a v e c   l a   f o r m a t i o n   d u n e   c o u c h e   i n t e r m t d i a i r e   d b x y c a r b u r e   T a C , O ,   s i t u & e   d   l ’ i n t e r f a c e   T a , C T a , O s .   D a n s   c e   c a s ,   i l   n ’ e s t   p a s   p o s s i b l e   d ’ e x c l u r e   u n   p r o c e s s u s   d t j h u s i o n n e l   l i m i t a n t   h   t r a v e r s   c e t t e   s o u s c o u c h e   d b x y n i t r u r e   d ’ e p a i s s e u r   c o n s t a n t e   a v e c   l e   t e m p s .   S o   f a r ,   v e r y   f e w   p a p e r s   a p p e a r   t o   h a v e   b e e n   p u b l i s h e d   o n   t h e   o x i d a t i o n   b e h a v i o u r   o f   t a n t a l u m   c a r b i d e   T a C   a n d   n o n e   o n   t h e   h e m i c a r b i d e   T a , C .   T h e   o x i d a t i o n   o f   h o t p r e s s e d   T a C   h a s   b e e n   s t u d i e d   b e t w e e n   1 7 0 0   a n d   2 2 O O ” C ,   f o r   a n   o x y g e n   p a r t i a l   p r e s s u r e   v a r y i n g   f r o m   1 0 m 3   t o   1 0 l   P a . ”   N o   o x i d e   l a y e r   w a s   f o r m e d ,   b u t   t h e   a u t h o r s   o b s e r v e d   a   c a r b o n   d i f f u s i o n   f r o m   t h e   h e a r t   t o   t h e   s u r f a c e   w h i c h   w a s   c o n s i d e r e d   a s   t h e   l i m i t i n g   s t e p .   T h e y   c o n c l u d e d ,   o n   t h e   b a s i s   o f   t h e   f e w   r e s u l t s   o b t a i n e d ,   t h a t   t h e   o x i d a t i o n   b e h a v i o u r   o f   T a C   a n d   N b C   w a s   q u a l i t a t i v e l y   t h e   s a m e .   C o n s e q u e n t l y ,   i n   v i e w   o f   t h i s   g r e a t   l a c k   o f   d a t a ,   i t   w a s   o f   i n t e r e s t   t o   i n v e s t i g a t e   t h e   d i f f e r e n c e   i n   r e a c t i v i t y   o f   t h e   t w o   t a n t a l u m   c a r b i d e s .   1 3 2 5   \\x0c', '1326   M. Desmaison-Brut   1   0.8   0.6   0.4   400   600   700   800   900   samples   X   x   again   Powders   4   tartice paramelers   a N.4   c fnm)   Meon  diameter   f&q (I4   Spe@r  area  (m2k)   Elements {mass %f   C   0   Fe   e t   a l .   2   E x p e r i m e n t a l   P r o c e d u r e   T h e   t a n t a l u m   c a r b i d e   p o w d e r ,   T a C ,   w a s   s u p p l i e d   b y   H .   C .   S t a r c k   a n d   t h e   c o m m e r c i a l l y   a v a i l a b l e   p o w d e r ,   T a $ Z ,   b y   C e r a c .   T h e   m a i n   c h a r a c t e r i s t i c s   o f   t h e   s t a r t i n g   p o w d e r s   a r e   p r e s e n t e d   i n   T a b l e   1 .   T h e   s a m e   H I P   t r e a t m e n t   h a s   b e e n   a p p l i e d   o n   t h e   t w o   i n i t i a l   p o w d e r s .   A f t e r   b e i n g   c o l d   i s o s t a t i c a l l y   p r e s s e d   a t   2 0 0   M P a ,   t h e   g r e e n   c o m p a c t   w a s   i n t r o   d u c e d   i n s i d e   a   t i t a n i u m   c o n t a i n e r .   A   c a r b o n   l a y e r   w a s   n e c e s s a r y   t o   p r e v e n t   a n y   c h e m i c a l   r e a c t i o n   b e t w e e n   t h e   c a p s u l e   a n d   t h e   c o m p o n e n t .   A f t e r   d e g a s s i n g   i n   v a c u u m   a t   6 0 0 ° C   f o r   1 0   h ,   t h e   c o n   t a i n e r   w a s   s e a l e d   a n d   H I P e d   a t   1   6 3 0 ° C ,   u n d e r   1 9 5   M P a   p r e s s u r e .   T h e   d w e l l   t i m e   l a s t e d   2   h .   T h e   r e l a t i v e   d e n s i t i e s   o f   t h e   s p e c i m e n s   t r e a t e d   b y   H I P ,   w e r e   g r e a t e r   t h a n   9 8 % .   C a l c u l a t i o n s   d o n e   o n   t h e   d e n s e   T a C   m a t e r i a l   a n d   b a s e d   o n   t h e   l a t t i c e   p a r a m e t e r   v a l u e   ( 0 . 4 4 5 4   n m )   i n d i c a t e d   t h a t   t h e   f e e   c a r b i d e   i s   s l i g h t l y   u n s t o i c h i o m e t r i c   ( T a C , . , , ) .   T h e   o x i d a t i o n   r e s i s t a n c e   w a s   t e s t e d   i n   a   d y n a m i c   f l o w   o f   p u r e   o x y g e n   ( 5 . 6   1 0 m 3   l i t r e / s ) ,   a t   a t m o   s p h e r i c   p r e s s u r e ,   u s i n g   a   S e t a r a m   m i c r o b a l a n c e .   C u b i c   s a m p l e s   ( 4   m m   s i d e )   w e r e   p o l i s h e d ,   w a s h e d   i n   a l c o h o l   w i t h   u l t r a s o n i c   a s s i s t a n c e   a n d   d r i e d .   T h e   p r o c e d u r e   w a s   a s   f o l l o w s :   f i r s t ,   t h e   f u r n a c e   w a s   e v a c u a t e d   ( l o 4   P a )   a n d   a   s t r e a m   o f   a r g o n   w a s   i n t r o d u c e d .   A s   t h e   t e m p e r a t u r e   w a s   i n c r e a s e d ,   t h e   s p e c i m e n   w a s   k e p t   o u t   o f   t h e   h o t   z o n e   u n t i l   1 5   m i n   a f t e r   t h e   i n t r o d u c t i o n   o f   o x y g e n   w h e n   i t   w a s   l o w e r e d   w i t h   a   m a g n e t i c   d e v i c e   i n t o   t h e   h o t   z o n e .   T h e   z e r o   t i m e   w a s   t a k e n   w h e n   t h e   p i a t i n u m   c r u   c i b l e   c o n t a i n i n g   t h e   s a m p l e   r e a c h e d   t h e   h o t   z o n e .   T h e   k i n e t i c   c u r v e s   w e r e   o b t a i n e d   b y   p l o t t i n g   t h e   f r a c t i o n a l   w e i g h t   c h a n g e   ( Y   ( a   =   A m / A m , )   v e r s u s   t i m e .   T h e   w e i g h t   g a i n   A m ,   c o r r e s p o n d s   t o   t h e   c o m p l e t e   o x i d a t i o n   e v a l u a t e d   b y   c o n s i d e r i n g   t h a t   T a C   o r   T a , C   i s   t r a n s f o r m e d   i n t o   T a * O ,   a c c o r d i n g   t o   t h e   r e a c t i o n s :   5 0 0   F i g .   1 .   N o n i s o t h e r m a l   o x i d a t i o n   c u r v e s   f o r   c u b i c   e x p o s e d   t o   f l o w i n g   o x y g e n   a t   1   a t m .   3   O x i d a t i o n   B e h a v i o u r   o f   T a n t a l u m   C a r b i d e ,   T a C   3 . 1   E f f e c t   o f   t e m p e r a t u r e   I s o t h e r m a l   c u r v e s   h a v e   b e e n   r e c o r d e d   a t   a t m o   s p h e r i c   p r e s s u r e ,   b e t w e e n   7 5 0   a n d   8 5 0 ° C   ( F i g .   2 ) .   T h e i r   s h a p e   i s   q u a s i l i n e a r   a t   l o w   t e m p e r a t u r e s   b u t   s h o w s   s o m e   d e c e l e r a t i o n   i n   t h e   u p p e r   t e m   p e r a t u r e   r a n g e .   A f t e r   b e i n g   k e p t   5   h   a t   S S O ” C ,   t h e   s p e c i m e n   i s   t o t a l l y   o x i d i z e d .   A   m a s t e r   r u n   w a s   c h o s e n   i n   t h e   m i d d l e   o f   t h e   s e r i e s   a n d   a   f a c t o r   A   w a s   c a l c u l a t e d   f o r   a n y   c u r v e   s u c h   t h a t   m u l t i p l i   c a t i o n   o f   t h e   t i m e   s c a l e   o f   t h e   r u n   b y   A   w o u l d   s u p e r i m p o s e   o n t o   t h e   m a s t e r   r u n   c u r v e .   E a c h   c u r v e   c a n   b e   s u p e r i m p o s e d   o n t o   a n y   o t h e r   b y   s u c h   a n   a f f i n i t y   r e l a t i o n s h i p   w i t h   t i m e   ( F i g .   3 ) .   A s   l o g   A   i s   f o u n d   t o   b e   a   l i n e a r   f u n c t i o n   o f   l / T ,   t h e   a c t i v a t i o n   e n e r g y   v a l u e   c a l c u l a t e d   ( E   =   3 7 9   f   1 6   k J / m o l )   i s   u n i q u e   o v e r   t h i s   r a n g e   o f   t e m p e r a t u r e .   3 . 2   E f f e c t   o f   p r e s s u r e   2   T a C   +   9 / 2   O 2   f   T a , O ,   +   2   C O ,   ( 1 )   T a , C   +   7 / Z   O 2   +   T a , 0 5   +   C O ,   ( 2 )   I n   o r d e r   t o   h a v e   a n   i d e a   o f   t h e   r e a c t i v i t y   o f   t h e   m a t e r i a l s ,   a   s a m p l e   o f   e a c h   t y p e   w a s   e x p o s e d   t o   o x y g e n   w i t h   a   l i n e a r   i n c r e a s i n g   t e m p e r a t u r e   o f   1 7   X   1 0 3 ” C / s .   T h e   n o n i s o t h e r m a l   o x i d a t i o n   c u r v e s   o f   F i g .   1   s h o w   t h a t   t h e   h e r n i c a r b i d e   m a t e r i a l   i s   m o r e   o x i d a t i o n   r e s i s t a n t   t h a n   t h e   m o n o c a r b i d e .   I s o b a r i c   c u r v e s   h a v e   b e e n   r e c o r d e d   a t   8 0 0 ° C   i n   t h e   p r e s s u r e   r a n g e   O 2   t o   1   1 0 ’   P a ,   a n d   p l o t t e d   a s   a   f u n c t i o n   o f   t i m e   ( F i g .   4 ) .   T h e   g e n e r a l   s h a p e   o f   t h e   k i n e t i c s   i s   r e t a i n e d   a n d   t h e   c u r v e s   c a n   b e   s u p e r i m p o s e d   ( F i g .   5 ) .   3 . 3   M o r p h o l o g i c a l   o b s e r v a t i o n s   D u r i n g   t h e   f i r s t   m i n u t e s   o f   t h e   r e a c t i o n ,   a   g r e y   o x i d e   f i l m   i s   f o r m e d .   A s   t i m e   a n d   t e m p e r a t u r e   i n c r e a s e ,   t h e   s i n t e r i n g   o f   t h e   o x i d e   s c a l e ,   f a c i l i t a t e d   b y   t h e   p r e s e n c e   o f   t h e   i n i t i a l   p o w d e r   i m p u r i t i e s ,   T a b l e   1 .   P o w d e r s   c h a r a c t e r i s t i c s   a n d   c o m p o s i t i o n   T a C   1 4 . 3 4   T a , C   1 4 . 9 0   0 4 4 5 4   O 3 1 0 3   0 . 4 9 3 7   2 . 0   4 . 5   1 . 1 6   0 . 2 5   6 . 2 9   0 . 2 6   0 . 0 0 4   3 . 1 7   0 . 2 5   0 . 0 0 1   \\x0c', \"Oxidation behaviour of HIPed   tantalum carbides   1327   0.8   0   10   0   a   i s   p e r f o r m e d .   F o r   a   g i v e n   t e m p e r a t u r e   ( 8 O O ” C ) ,   t h e   o x i d e   s c a l e   t h i c k n e s s   m a y   b e   r e l a t e d   t o   t h e   r e a c t i o n   t i m e   b y   a   l i n e a r   l a w   a n d   t h e   t i m e   d e p e n   d e n c e   o f   t h e   T a C   ( u n r e a c t e d )   c o r e   t h i c k n e s s   i s   g i v e n   i n   F i g .   6 .   X r a y   a n a l y s i s   s h o w s   o n l y   t h e   p r e s e n c e   o f   t h e   o x i d e   j 3 T a , 0 5   i n d e p e n d e n t l y   o f   p r e s s u r e   a n d   t e m p e r a t u r e .   F i g u r e s   7 ( a )   a n d   7 ( b )   s h o w   t h e   c o l u m n a r   a s p e c t   o f   t h e   t a n t a l u m   o x i d e .   T h e   h i g h   v a l u e   o f   t h e   P i l l i n g   a n d   B e d w o r t h   r a t i o   ( A   =   2 . 1 3 )   i n d u c e s   s t r e s s e s   i n   t h e   s c a l e .   T h e   g o o d   a d h e r e n c e   o f   t h e   o x i d e   t o   t h e   c o r e   d o e s   n o t   a l l o w   s t r e s s   r e l a x a t i o n   b y   s p a l l i n g   b u t   l e a d s   t o   t h e   o p e n   i n g   o f   t h e   c u b e   e d g e s   a n d   t o   t h e   f o r m a t i o n   o f   a   m a l t e s e   c r o s s   ( F i g .   7 ( b ) ) .   T h e   s a m e   p r o c e s s   h a s   a l r e a d y   b e e n   o b s e r v e d   i n   t h e   c a s e   o f   t h e   o x i d a t i o n   b e h a v i o u r   o f   a   d e n s e   a n d   H I P e d   T a N   m a t e r i a l . 1 2   T h e   o x i d e   s u r f a c e   a s p e c t   i s   d i f f e r e n t   a t   8 2 5 ° C   o r   l o w e r   t h a n   a t   8 5 0 ° C   ( F i g s   7 ( c )   a n d   ( d ) )   c e r t a i n l y   d u e   t o   t h e   b e t t e r   s i n t e r i n g   o f   t h e   e x t e r n a l   p a r t .   3 . 4   I n t e r p r e t a t i o n   T h e   k i n e t i c   r e s u l t s   i n d i c a t e   t h a t   t h e   r e a c t i o n   r a t e   V   i s   a   f u n c t i o n   o f   p r e s s u r e   a n d   t e m p e r a t u r e .   A s   b o t h   i s o t h e r m a l   a n d   i s o b a r i c   c u r v e s   a r e   i n   a   c l o s e   1   0 . 6   0 . 4   0 . 2   0   5   1 5   2 0   F i g .   2 .   I s o t h e r m a l   c u r v e s   r e c o r d e d   a t   1   a t m   i n   t h e   t e m p e r a t u r e   r a n g e   7 % 8 5 0 ° C .   1   0 . 8   0 . 6   0 . 4   0 . 2   Y .   t i m e   0 1 )   5   1 0   1 5   F i g .   3 .   S u p e r i m p o s e d   i s o t h e r m a l   c u r v e s .   2 0   a f f i n i t y   r e l a t i o n s h i p   w i t h   t i m e ,   t h e   e q u a t i o n   c a n   b e   w r i t t e n   i n   a   s e p a r a t e d   v a r i a b l e s   f o r m :   V =   $   =   f ( a ) g ( T ) h ( P )   =   C t e   f ( a ) h ( P ) e x p   &   ( 3 )   w h e r e   f ( a )   i s   a   m o r p h o l o g i c a l   t e r m   c h a r a c t e r i s   t i c   o f   t h e   r e a c t i o n   a r e a   a n d   C t e   i s   a   c o n s t a n t .   1   0 . 8   0 . 6   0 . 4   0 . 2   0   F i g .   1   0 . 8   0 . 6   0 . 4   0 . 2   0   3   2 ”   t i m e   ( h )   0   5   1 0   1 5   2 0   4 .   I s o b a r i c   c u r v e s   r e c o r d e d   a t   8 0 0 ° C   i n   t h e   p r e s s u r e   r a n g e   0 . 2   t o   1   X   1 0 ”   P a .   r   I   *   0 . 6   1 0   % ' a   I   5   1 0   1 5   2 0   F i g .   5 .   S u p e r i m p o s e d   i s o b a r i c   c u r v e s .   c u b e   e d g e   ( m m )   0   2   4   6   a   1 0   1 2   F i g .   6 .   T i m e   d e p e n d e n c e   o f   t a n t a l u m   c a r b i d e   c o r e   t h i c k n e s s   a t   8 0 0 ° C .   t i m e   ( I I )   \\x0c\", '1328   M. Desmaison-Bru?   e t   a l .   6 4   ( d )   F i g .   7 .   S E M   o f   o x i d i z e d   T a C   m a t e r i a l s :   ( a )   8 O O ” C ,   1 5   m i n .   c r o s s s e c t i o n ;   ( b )   8 O O ” C ,   4   h ,   c r o s s s e c t i o n ;   ( c )   8 2 5 ° C .   2 0   h ,   s u r f a c e ;   ( d )   8 5 O ” C ,   2 0   h ,   s u r f a c e .   T h i s   m e a n s   t h a t   t h e   c o n t r o l l i n g   p r o c e s s   i s   u n i q u e   a n d   i d e n t i c a l   t o   i t s e l f   o v e r   t h e   w h o l e   r e a c t i o n   r a n g e . ‘ 2 . ‘ ”   A t   c o n s t a n t   t e m p e r a t u r e   a n d   p r e s s u r e ,   t h e   s h a p e   o f   t h e   c u r v e s   i s   d e t e r m i n e d   o n l y   b y   g e o   m e t r i c a l   f a c t o r s .   i n n e r   i n t e r f a c e .   T h e r e f o r e ,   t h e   r e a c t i o n   r a t e   V   i s   d i r e c t l y   p r o p o r t i o n a l   t o   t h e   s u r f a c e   a r e a   o f   t h e   n o n o x i d i z e d   c u b i c   c a r b i d e   c o r e   o f   e d g e   a ,   a t   t i m e   t .   T h e   o x i d e   p o r o s i t y   a n d   t h e   o p e n i n g   o f   t h e   c u b e   e d g e s   a t   t h e   b e g i n n i n g   o f   t h e   r e a c t i o n   a l l o w   d i r e c t   a c c e s s   o f   o x y g e n   a t   t h e   c a r b i d e o x i d e   i n t e r f a c e .   O n   t h e   o t h e r   h a n d ,   t h e   c a r b o n   d i o x   i d e   g a s   r e l e a s e   f r o m   t h e   c o r e   i s   e a s y .   A l l   t h e s e   r e m a r k s   s u g g e s t   a   r e a c t i o n   r e g i m e   a t   t h i s   V   =   $   =   K ( T , P ) S ( ( w )   ( 4 )   S ( a )   =   6   a 2   b e i n g   r e l a t e d   t o   t h e   c o n v e r s i o n   d e g r e e   L Y   t h r o u g h   a .   I n d e e d :   c x   =   1     V / V ,   \\x0c', 'K2PO2   850°C   0   Fig.   8.   0   Fig.   9.   0.8   0   Fig.   10.   0.04   J/mole   10%   (l/l   0   0.2   Fig.   11.   O x i d a t i o n   b e h a v i o u r   o f   H I P e d   t a n t a l u m   c a r b i d e s   1 3 2 9   w i t h   Y   t h e   v o l u m e   o f   t h e   r e s i d u a l   s u b s t r a t e   a n d   v .   t h e   v o l u m e   o f   t h e   i n i t i a l   c u b e   o f   e d g e   a , .   B y   e x t e n s i o n :   T h e   c o n s t a n t   d e p e n d e n c e   w i t h   p r e s s u r e   m a y   b e   r e p r e s e n t e d   b y   a   r e l a t i o n   o f   t h e   f o r m :   ( 5 )   a   =   a ,   ( 1     , ) l n   T h e   r a t e   l a w   b e c o m e s :   ( 6 )   d o l !   d t   =   6   K ( T , P )   a t   ( 1     a ) 2 1 3   I n t e g r a t i o n   o f   e x p r e s s i o n   ( 7 )   l e a d s   t o :   F ( a )   =   1     ( 1     C Y ) “ ~   =   k t   w i t h   k   =   2   a i   K ( T , P ) .   ( 8 )   T h e r e f o r e ,   t h e   a s s e s s m e n t   o f   t h e   v a l i d i t y   o f   o u r   k i n e t i c   m o d e l   m a y   b e   f o u n d   b y   p l o t t i n g   F ( a )   v e r s u s   t   ( F i g .   8 ) .   A s   c a n   b e   s e e n ,   l i n e a r   p l o t s   a r e   o b t a i n e d   f o r   a l l   t e m p e r a t u r e s .   B y   p l o t t i n g   l o g   k   v e r s u s   r e c i p r o c a l   t e m p e r a t u r e   ( F i g .   9 )   a   v a l u e   o f   3 8 5   _ +   1 1   k J / m o l e   i s   c a l c u l a t e d .   I n   t h e   s a m e   w a y ,   a t   v a r i o u s   p r e s s u r e s ,   a   l i n e a r   r e l a t i o n   i s   o b t a i n e d ,   f o r   e x a m p l e   a t   8 0 0 ° C   b y   p l o t t i n g   F ( a )   v e r s u s   t i m e   ( F i g .   1 0 ) .   1   0 . 8   ‘ ( a )   a   0 . 6   0 . 4   0 . 2   0   t o . 6   0 . 4   I   0 . 2   5   1 0   1 5   2 0   O x i d a t i o n   c u r v e s   p l o t t e d   a g a i n s t   t i m e   i n   t h e   c u b i c   f o r m ,   a t   v a r i o u s   t e m p e r a t u r e s .   1   L n   ( I Q   1   A c t i v a t i o n e n e r g y   E r 3 8 5 f l l k   2   3   4   5   6   8 . 8   9 . 2   9 . 6   1 0   I n f l u e n c e   o f   t e m p e r a t u r e   o n   t h e   r a t e   c o n s t a n t   K   a t   1   a t m .   k = a P l ( l   + b P )   ( 9 )   w h e r e :   a   =   0 . 0 6 7   a n d   b   =   0 . 7 8 1   ( F i g .   1 1 )   T h e   o b s e r v e d   p r e s s u r e   d e p e n d e n c e   m a y   b e   e x p l a i n e d   b y   a s s u m i n g   t h a t   a n   a d s o r p t i o n   e q u i l i b r i u m   o f   m o l   e c u l a r   o x y g e n   p r e c e e d s   t h e   f o r m a t i o n   o f   T a 2 0 , :   0 2 + s e o 2 s   ( 1 0 )   A c c o r d i n g   t o   t h e   L a n g m u i r   m o d e l ,   t h e   f r a c t i o n   0 ,   o f   a d s o r p t i o n   s i t e s   s   o c c u p i e d   b y   c h e m i s o r b e d   o x y g e n   i s :   0 ,   =   K J b   1   +   K , P ,   2   ( 1 1 )   w h e r e   K ,   i s   t h e   e q u i l i b r i u m   c o n s t a n t   o f   r e a c t i o n   ( 1 0 ) .   I f   t h e   o x i d a t i o n   r a t e   i s   p r o p o r t i o n a l   t o   t h e   n u m b e r   o f   a d s o r b e d   o x y g e n   m o l e c u l e s ,   v   =   k o t I G   =   1   +   K , P o   2   ( 1 2 )   w i t h   K 2   =   k o K ,   1   0 . 8   F ( a )   / d a , l   a   1   0 . 6   0 . 4   0 . 2   0   5   1 0   1 5   2 0   O x i d a t i o n   c u r v e s   D l o t t e d   a g a i n s t   t i m e   i n   t h e   c u b i c   f o r m ,   a t   v a r i o b s   p r e s s u r e s   ( S O O T )   0 . 0 3   0 . 0 2   0 . 0 1   0   0 . 4   0 . 6   0 . 8   1   E f f e c t   o f   o x y g e n   p r e s s u r e   o n   t h e   r a t e   c o n s t a n t   a t   8 0 0 ° C .   \\x0c', '1330   0.8   0.6   0.4   0.2   0   0   Fig.   5   10   I5   20   time (h)   0   Fig.   10   15   20   Ln   energy   10   a   10   15   20   Fig.   15.   0.8   0   Fig.   16.   0.06   0   Fig.   10   15   0   02   Fig.   17.   1   M .   D e s m a i s o n B r u t   e t   a l .   F i n a l l y ,   a   p a r a m e t r i c   k i n e t i c   l a w   o f   t y p e   ( 3 )   m a y   4   O x i d a t i o n   B e h z w i o u r   o f   T a n t a l u m   H e m i c a r b i d e   r e p r e s e n t   t h e   o x i d a t i o n   r a t e :   T a 2 C   1   V   =   C t e   ( 1     ( Y ) * ‘ ~   0 . 0 6 7   P / ( 1   +   0 . 7 8 1   P )   e x p ( 3 8 5   O O O / R T )   ( 1 3 )   1 2 .   I s o t h e r m a l   c u r v e s   r e c o r d e d   a t   I   a t m   i n   t h e   t e m p e r a t u r e   r a n g e   7 5 0 8 5 0 ° C .   u   t i m e   ( h )   ,   I   5   5   1 3 .   I s o b a r i c   c u r v e s   r e c o r d e d   a t   8 O O ” C ,   i n   t h e   p r e s s u r e   r a n g e   0 . 2   t o   1   P a .   O x i d a t i o n   c u r v e s   p l o t t e d   a g a i n s t   t i m e   i n   t h e   c u b i c   X   l o 5   f o r m ,   a t   v a r i o u s   p r e s s u r e s .   0 . 6   .   5   2 0   D A   0 . 6   0 . 8   1 4 .   O x i d a t i o n   c u r v e s   p l o t t e d   a g a i n s t   t i m e   i n   t h e   c u b i c   E f f e c t   o f   o x y g e n   p r e s s u r e   o n   t h e   r a t e   c o n s t a n t   a t   f o r m ,   a t   v a r i o u s   t e m p e r a t u r e s .   S O O T .   4 . 1   E f f e c t   o f   t e m p e r a t u r e   a n d   p r e s s u r e   A n   i d e n t i c a l   s t u d y   h a s   b e e n   p e r f o r m e d   o n   t h e   t a n   t a l u m   h e m i c a r b i d e   c e r a m i c .   T h e   r e s u l t s   h a v e   b e e n   2   ( k )   A c t i v a t i o n   E   = 1 2 9 f 7   l & M o l e   3   4   1 0 4 / T ( 1 / K )   5   8 . 8   9 . 2   9 . 6   I n f l u e n c e   o f   t e m p e r a t u r e   o n   t h e   r a t e   c o n s t a n t   K   a t   1   0 . 6   1   a t m .   K   0 . 0 5     K   _   w 8 3 . P   1 + 0 , 9 2 9 . P   0 . 0 4   =   P r e s a i o n ( l ~ P a )   \\x0c', 'Oxidation behaviour of HIPed   tantalum carbides   1331   Morphological observations   Fig. 18.   a n a l y s e d   b y   t h e   s a m e   m a t h e m a t i c a l   a p p r o a c h   a n d   a p p e a r   i n   F i g s   1 2   1 7 .   T h e   g l o b a l   k i n e t i c   l a w   m a y   b e   w r i t t e n   i n   t h e   f o r m :   V   =   C t e ( 1     a ) 2 ’ 3   O 0 8 3   P / ( 1   +   0 . 9 2 9   P )   e x p (   1 2 9   0 0 0 / R   r )   ( 1 4 )   4 . 2   X r a y   a n a l y s i s   s h o w s   t h e   p r e s e n c e   o f   p   T a 2 0 S   i n d e p e n d e n t l y   o f   p r e s s u r e   a n d   t e m p e r a t u r e .   T h e   h e m i p e n t o x i d e   w h i c h   f o r m s   a   m a l t e s e   c r o s s   i s   r a t h e r   d e n s e   a n d   a d h e r e n t   t o   t h e   c o r e   ( F i g s   1 8 ( a )   a n d   ( b ) ) .   B y   o p t i c a l   m i c r o s c o p y ,   w e   h a v e   o b s e r v e d   t h a t   t h e   o x i d a t i o n   s t a r t s   a t   t h e   g r a i n   b o u n d a r i e s   o f   t h e   T a 2 C   c e r a m i c .   O n   c r o s s s e c t i o n s ,   a   l a y e r   w i t h   a   m a x i m u m   t h i c k n e s s   o f   2 0 2 5   k m   i s   p r e s e n t   a t   t h e   i n t e r f a c e   T a 2 0 S T a , C   ( F i g .   1 8 ( c ) ) .   T h e   m a x   i m u m   t h i c k n e s s   s e e m s   t o   b e   i n d e p e n d e n t   o f   t i m e   b u t   t h e   l a y e r   d i s a p p e a r s   n e a r   t h e   e d g e s   o f   t h e   c u b e .   T h e   m i c r o h a r d n e s s   v a l u e   o f   t h e   T a , O ,   p h a s e   i s   l o w   ( H ,   5 N   3 . 5   +   0 . 5   G P a )   c o m p a r e d   t o   t h e   m i c r o h a r d n e s s e s   o f   t h e   i n t e r m e d i a t e   p h a s e   ( a )   0 4   M i c r o g r a p h s   o f   o x i d i z e d   T a $   m a t e r i a l s :   ( a )   7 5 0 ° C   2 0   h   ( o p t i c a l   m i c r o g r a p h ) ;   ( b )   7 5 0 ° C   2 0   h   ( s c a n n i n g   e   m i c r o g r a p h ) ;   ( c )   8 O O ” C ,   2 0   h   ( o p t i c a l   m i c r o g r a p h ) .   : l e c l   i r o n   \\x0c', '(12-l   M. Desmaison-Brut   et ~1.‘~   1 3 3 2   e t   a l .   f   0 . 9   G P a )   a n d   o f   t h e   t a n t a l u m   h e m i c a r b i d e   T a &   ( 1 0 . 7   f   0 . 9   G P a ) .   B y   X r a y   a n a l y s i s   o f   s u c c e s s i v e l y   p o l i s h e d   s u r f a c e s   t h r o u g h   t h e   i n t e r m e d i a t e   l a y e r ,   w e   h a v e   n o t i c e d   a n   e v o l u t i o n   o f   t h e   l a t t i c e   p a r a m e t e r s   o f   t h e   T a &   g r a d e .   T h e   E D A X   a n a l y s i s   o f   t h e   c a r b o n ,   o x y g e n   a n d   t a n t a l u m   e l e m e n t s   a n d   t h e i r   E P M A   p r o f i l e s   c l e a r l y   s h o w   t h e   p r e s e n c e ,   t h r o u g h   t h e   i n t e r f a c e ,   o f   b o t h   c a r b o n   a n d   o x y g e n 1 4   s u g g e s t i n g   t h e   f o r m a t i o n   o f   a   t a n t a l u m   c a r b o x i d e   l a y e r .   5   O x i d a t i o n   i n   T h e s e   a n d   O t h e r   C a r b i d e   S y s t e m s   5 . 1   C o m p a r i s o n   o f   t h e   o x i d a t i o n   b e h a v i o u r   o f   t h e   t w o   c a r b i d e s   T a C   a n d   T a 2 C   I n   o u r   c a s e ,   t h e   c o r r e l a t i o n   o f   m o r p h o l o g i c a l   o b s e r v a t i o n s   w i t h   k i n e t i c   r e s u l t s   i s   r e a d i l y   m a d e .   I n   b o t h   c a s e s ,   t h e   d e n s i f i c a t i o n   o f   t a n t a l u m   p e n   t o x i d e   w a s   n o t   s u f f i c i e n t   t o   l e a d   t o   a   r e s t r i c t e d   a c c e s s   o f   o x y g e n   t o   t h e   T a C   s u r f a c e ;   a s   a   c o n s e   q u e n c e ,   t h e r e   w a s   n o   r e q u i r e m e n t   f o r   t h e   e x i s t e n c e   o f   a   d i f f u s i o n l i m i t i n g   s t e p .   I n   a d d i t i o n ,   d u e   t o   t h e   s a m p l e   g e o m e t r y ,   a   m a l   t e s e   c r o s s   i s   f o r m e d .   H o w e v e r ,   t h e   T a &   c o r e   d o e s   n o t   k e e p   a s   p e r f e c t   a   c u b i c   s h a p e   a s   i t   d o e s   i n   t h e   c a s e   o f   t h e   o x i d a t i o n   o f   t a n t a l u m   c a r b i d e   T a C   ( F i g s   7   a n d   1 8 ( a ) ) .   R a m a n   s p e c t r a   i n   t h e   l a t t e r   c a s e   h a v e   n o t   s u g g e s t e d   t h e   p r e s e n c e   o f   f r e e   c a r b o n   w i t h i n   t h e   s c a l e .   T h e r e f o r e ,   f o r   T a C ,   t h e   l i b e r a t e d   c a r b o n   m a y   r e a c t   d i r e c t l y   w i t h   o x y g e n   t o   f o r m   C O z .   T h e   c o l u m n a r   o x i d e   s c a l e   a l l o w s   a n   e a s y   o u t w a r d   d i f f u s i o n   o f   t h i s   g a s .   I n   b o t h   c a s e s ,   t h e   r e a c t i o n s   a p p e a r   t o   b e   g o v   e r n e d   b y   a n   i n t e r f a c i a l   p r o c e s s 2 1 2 3   l o c a t e d   a t   t h e   c a r b i d e o x i d e   i n t e r f a c e .   H o w e v e r ,   t h e   h e m i c a r b i d e   e x h i b i t s   a   h i g h e r   o x i d a t i o n   r e s i s t a n c e   p e r h a p s   d u e   t o   t h e   t r a n s i t o r y   f o r m a t i o n   o f   t h e   o x y c a r b i d e   p h a s e .   T h e   m o n o c a r b i d e   T a C   h a s   a   f e e   s t r u c t u r e   w h e r e   t h e   c a r b o n   a t o m s   a r e   l o c a t e d   i n s i d e   t h e   o c t a h e d r a l   s i t e s .   W h e n   t h e   m a t e r i a l   i s   u n s t o i c h i o m e t r i c ,   t h e   v a c a n t   s i t e s   c a n   b e   o c c u p i e d   b y   o x y g e n   a t o m s   b u t   h e r e   a   l i m i t e d   n u m b e r   o f   s u c h   s i t e s   a r e   a v a i l a b l e   a s   t h e   m o n o c a r b i d e   i s   a l m o s t   s t o i c h i o m e t r i c   ( T a C , , . & .   T h e   p r e s e n c e   o f   a n   o x y c a r b i d e   l a y e r   h a s   n o t   b e e n   d e t e c t e d   e v e n   a f t e r   a   f e w   m i n u t e s   o f   o x i   d a t i o n   t r e a t m e n t .   O n   t h e   o t h e r   h a n d ,   t h e   h e x a g o n a l   a n t i C d I ,   t y p e   c r y s t a l   s t r u c t u r e   o f   T a , C   w h e r e   t h e   c a r b o n   a t o m s   a r e   o c c u p y i n g   o n l y   o n e   p l a n e   o u t   o f   t w o ,   c o u l d   f a c i l i t a t e   t h e   d i s s o l u t i o n   o f   o x y g e n   a n d   t h e   i n s e r t i o n   o f   o x y g e n   a t o m s   i n s i d e   t h e   T a , C   n e t   w o r k   w i t h   t h e   f o r m a t i o n   o f   a   t a n t a l u m   c a r b o x i d e   l a y e r   p r e c e d i n g   t h e   g r o w t h   o f   t h e   h e m i p e n t o x i d e   s c a l e .   T h e   o b s e r v a t i o n   o f   c r o s s s e c t i o n s   s h o w s   t h a t   t h i s   r e a c t i o n   s t a r t s   p r e f e r e n t i a l l y   a t   t h e   g r a i n   b o u n   d a r i e s   o f   t h e   s u b s t r a t e   a n d   p r o c e e d s   a t   a n   a l m o s t   c o n s t a n t   s p e e d .   T h e r e f o r e ,   i n   t h i s   c a s e ,   w e   c a n n o t   e x c l u d e   a   d i f f u s i o n   l i m i t i n g   s t e p   t h r o u g h   a n   o x y   n i t r i d e   s u b l a y e r   o f   c o n s t a n t   t h i c k n e s s   w i t h   t i m e .   5 . 2   C o m p a r i s o n   w i t h   o t h e r   p u b l i s h e d   w o r k s   o n   o x i d a t i o n   o f   c a r b i d e   m a t e r i a l s   T h e   f o r m a t i o n   o f   s u c h   i n t e r l a y e r s   h a s   b e e n   r e p o r t e d   i n   o t h e r   s t u d i e s   d e v o t e d   t o   t h e   o x i d a t i o n   o f   s e v e r a l   t r a n s i t i o n   c a r b i d e s   ( H f C r 5 ‘ * ,   Z r C ’ 9 2 3 ,   T i C 2 4 ,   N b C 2 5 ) .   A f t e r   h a f n i u m   c a r b i d e   C V D   f i l m s   h a v e   b e e n   o x i d i z e d   a t   t e m p e r a t u r e s   i n   t h e   r a n g e   o f   1 4 0 0   t o   2 0 6 O ” C ,   t h r e e   d i s t i n c t   l a y e r s   a r e   o b s e r v e d :   ( a )   a   r e s i d u a l   h a f n i u m   c a r b i d e   l a y e r ,   ( b )   a   d e n s e   a p p e a r i n g   h a f n i u m   o x i d e   i n t e r l a y e r   c o n t a i n i n g   c a r b o n ,   a n d   ( c )   a   p o r o u s   o u t e r   l a y e r   o f   h a f n i u m   o x i d e . 1 5   A s   t h e   o v e r a l l   d u p l e x   o x i d i z e d   l a y e r   a p p e a r e d   t o   b e   p r o t e c t i v e ,   B e r g e r o n   s o l v e d   t h e   d i f f u s i o n a l   p r o b l e m   b y   c o m b i n i n g   e x p e r i m e n   t a l   m e a s u r e m e n t s   o f   l a y e r   t h i c k n e s s e s   a n d   o x y g e n   c o n c e n t r a t i o n s   w i t h   a n   e x t e n d e d   f o r m u l a t i o n   o f   t h e   m o v i n g b o u n d a r y   d i f f u s i o n   t h e o r y .   T h e   r e s u l t s   i n d i c a t e   t h a t   t h e   o x i d e   i n t e r l a y e r   i s   a   b e t t e r   d i f f u   s i o n   b a r r i e r   f o r   o x y g e n   t h a n   e i t h e r   o f   t h e   o t h e r   l a y e r s .   I n   t h e   c a s e   o f   n o t   f u l l y   d e n s i f i e d   s i n t e r e d   m a t e   r i a l s ,   t h e   f o r m a t i o n   o f   a   l e s s   p r o m i n e n t   d a r k   b l a c k   i n t e r l a y e r   h a s   a l s o   b e e n   d e s c r i b e d . 1 6   H o w e v e r ,   t h e   c a r b i d e   c o u l d   f i r s t   d i s s o l v e   o x y g e n ,   u p   t o   3 0 %   a t   2 0 0 0 ° C , 1 7   b e f o r e   i t   c o n v e r t s   t o   a   c o m p a c t   o x y g e n   d e f i c i e n t   o x i d e .   T h e   t r a n s f o r m a t i o n   o c c u r s   p r i o r   t o   f u l l   o x i d a t i o n   o f   t h e   c a r b i d e   w h e n   t h e r e   i s   i n s u f f i c i e n t   o x y g e n   t o   o x i d i z e   a l l   t h e   c a r b o n   t o   C O .   T h e   i n t e g r i t y   a n d   l o w   p o r o s i t y   o f   t h e   i n t e r   l a y e r   i n d i c a t e   t h a t   p o t e n t i a l   g a s   p r o d u c t i o n   i s   n o t   e f f e c t i v e   i n   c r e a t i n g   a   d i s r u p t i o n   a t   i n t e r f a c e s   o r   s e p a r a t i o n   b e t w e e n   l a y e r s . i 5   T h e   i s o t h e r m a l   o x i d a t i o n   o f   H f C   p o w d e r s ,   a t   r e l a t i v e l y   l o w   t e m p e r a t u r e s   ( 4 8 0 6 O O ” C ) ,   p r o c e e d s   b y   s i m i l a r   p r o c e s s e s .   I 8   A f t e r   a n   i n i t i a l   r a p i d   o x i d a   t i o n   w i t h   t h e   f o r m a t i o n   o f   o x y c a r b i d e ,   H f C , O , , ,   o x i d a t i o n   p r o c e e d s   b y   a   d i f f u s i o n c o n t r o l l e d   p r o   c e s s   i n   t h e   e a r l y   s t a g e   ( I O S O %   o x i d a t i o n )   f o l   l o w e d   b y   a   p h a s e b o u n d a r y c o n t r o l l e d   p r o c e s s   i n   t h e   l a t e r   s t a g e   ( ~ 5 0 % ) .   T h e   c h a n g e   i n   o x i d a t i o n   m e c h a n i s m   i s   a s s o c i a t e d   w i t h   t h e   g e n e r a t i o n   o f   c r a c k s   o n   t h e   g r a i n s ,   r e s u l t i n g   f r o m   t h e   g r o w t h   o r   e x p a n s i o n   s t r e s s   d u e   t o   t h e   f o r m a t i o n   o f   m o n o   c l i n i c   H f 0 2 .   T h e   f o r m a t i o n   o f   a m o r p h o u s   c a r b o n   i n   t h e   o x i d i z e d   s a m p l e s   w a s   s u g g e s t e d   b y   R a m a n   s p e c t r a ,   t h i s   c a r b o n   b e i n g   d i f f i c u l t   t o   r e m o v e   a t   l o w   o x y g e n   p r e s s u r e s   a n d   t e m p e r a t u r e s .   C a r b o n c o n t a i n i n g   o x i d e   s c a l e s   h a v e   a l s o   b e e n   o b t a i n e d   b y   o x i d a t i o n   o f   p o w d e r s , 1 9 , 2 0   s i n t e r e d   \\x0c', 'O x i d a t i o n   b e h a v i o u r   o f   H I P e d   t a n t a l u m   c a r b i d e s   1 3 3 3   m a t e r i a l s 2 1 * 2 2   o r   s i n g l e   c r y s t a l s 2 3   o f   z i r c o n i u m   c a r b i d e .   D u f o u r   e t   a L I 9   a n d   S h i m a d a   e t   a L 2 0   h a v e   s u g g e s t e d   t h e   f o r m a t i o n   o f   c a r b o n   d u r i n g   o x i d a   t i o n   o f   Z r C   p o w d e r s .   N o   s a t i s f a c t o r y   e x p l a n a t i o n   f o r   t h e   o x i d a t i o n   o f   Z r C   p o w d e r s   h a s   b e e n   g i v e n   i n   v i e w   o f   t h e   v a r i e d   k i n e t i c   r e s u l t s ,   a n d   w i t h   r e g a r d   t o   t h e   i n f l u e n c e   o f   c r a c k i n g   o n   o x i d a t i o n .   B a r n i e r   e t   a 1 . 2 ’ $ 2 2   a l s o   a s s u m e d   t h a t   a   c e r t a i n   a m o u n t   o f   c a r b o n   i s   r e t a i n e d   i n   t h e   o x y c a r b i d e   s t r u c t u r e .   T h e y   h a v e   r e p o r t e d   t h a t   t h e   o x i d a t i o n   k i n e t i c s   o f   s i n t e r e d   Z r C ,   i n   t h e   t e m p e r a t u r e   r a n g e   o f   4 0 0   t o   7 0 0 ° C ,   u n d e r   1 3 0   k P a   o f   o x y g e n ,   a r e   s u c c e s s i v e l y   c o n t r o l l e d   b y   a   d i f f u s i o n   p r o c e s s   a n d   b y   a   p h a s e b o u n d a r y   p r o c e s s .   M o r e   r e c e n t l y ,   S h i m a d a   e t   a L 2 ’   c a r r i e d   o u t   t h e   i s o t h e r m a l   o x i d a t i o n   o f   Z r C   s i n g l e   c r y s t a l s   w i t h   ( 1 0 0 )   o r i e n t a t i o n   a t   t e m p e r a t u r e s   o f   5 0 0 ” ,   5 5 0 ”   a n d   6 0 0 ° C   a t   a n   o x y g e n   p r e s s u r e   o f   2 . 6   k P a   f o r   t i m e s   u p   t o   2 4 0   h .   I t   w a s   f o u n d   t h a t   t h e   o x i d e   s c a l e   w a s   d i v i d e d   i n t o   t w o   r e g i o n s ,   z o n e s   1   a n d   2 ,   w h i c h   c o n t a i n e d   1 4 2 3   a n d   7 1 0   a t %   c a r b o n ,   r e s p e c t i v e l y .   T h e   t h i c k n e s s   o f   z o n e   1 ,   c o n s i s t i n g   o f   a   c o m p a c t ,   p o r e f r e e   m a t r i x   o f   c u b i c   Z r O , ,   i n c r e a s e d   p a r a b o l i c a l l y   u p   t o   2 4 0   h   a t   5 0 0 ° C   a n d   p r o b a b l y   i n   a n   e a r l y   p e r i o d   a t   5 5 0   a n d   6 0 0 ° C   r e a c h i n g   a   c o n s t a n t   t h i c k n e s s   o f   a b o u t   2 3   p m .   I n   c o n t r a s t ,   t h e   t h i c k n e s s   o f   z o n e   2 ,   w h e r e   s o m e   g r o w t h   a n d   a g g r e g a t i o n   o f   t h e   c Z r 0 2   o c c u r r e d ,   p r o d u c i n g   5   t o   2 0 n m   s i z e d   p a r t i c l e s ,   i n c r e a s e d   l i n e a r l y   w i t h   t i m e .   T h e   o x i d a t i o n   m e c h a n i s m   o f   t i t a n i u m   c a r b i d e   p o w d e r s   a t   l o w   t e m p e r a t u r e s   ( 3 5 0 5 0 0 ° C )   h a s   n o t   y e t   b e e n   c l e a r l y   e l u c i d a t e d .   S h i m a d a   e t   a 1 . 2 4   r e p o r t e d   t h e   f o r m a t i o n   o f   o x y c a r b i d e / t i t a n i u m   s u b o x i d e s   a n d   t h e   c r y s t a l l i z a t i o n   o f   a n a t a s e ,   f o l   l o w e d   b y   t h e   g e n e r a t i o n   o f   c r a c k s   i n   t h e   g r a i n s .   T h e   o x i d a t i o n   b e h a v i o u r   o f   p o w d e r   a n d   s i n g l e   c r y s t a l s   o f   n i o b i u m   c a r b i d e   i s   a l s o   c o m p l e x .   H o w   e v e r ,   r e c e n t   k i n e t i c   r e s u l t s 2 5   s u g g e s t e d   t h a t   t h e   o x i   d a t i o n   o f   b o t h   s a m p l e s   p r o c e e d s   e s s e n t i a l l y   b y   a   p h a s e b o u n d a r y c o n t r o l l e d   r e a c t i o n   d u e   t o   t h e   f o r m a t i o n   o f   c o l u m n a r ,   p o r o u s   N b 2 0 5   g r a i n s ,   w i t h   t h e   m a j o r   a x i s   n o r m a l   t o   t h e   s u r f a c e   f a c i l i t a t   i n g   C O 2   e v o l u t i o n .   6   M o d e l   f o r   t h e   O x i d a t i o n   M e c h a n i s m   o f   T a 2 C   N o n e   o f   t h e   p r e c e d i n g   a p p r o a c h e s   i s   a d e q u a t e   t o   e x p l a i n   t h e   p r e s e n t   r e s u l t s .   O n   t h e   c o n t r a r y ,   t h e   p a r a l i n e a r   o x i d a t i o n   m o d e 1 , 2 6   f i r s t   d i s c u s s e d   b y   L o r i e r s 2 7   i n   a n   i n t e r p r e t a t i o n   o f   t h e   o x i d a t i o n   o f   c e r i u m   ( g r o w t h   o f   p o r o u s   C e 0 2   o n   c o m p a c t   C e , O , )   c o u l d   b e   o f   p a r t i c u l a r   i n t e r e s t .   I n d e e d ,   t h i s   m o d e l ,   u s e d   b y   o t h e r   a u t h o r s ,   f o r   i n s t a n c e   f o r   t h e   o x i d a t i o n   o f   t u n g s t e n 2 8   a n d   c a l c i u m , 2 9   d e s c r i b e s   t h e   g r o w t h   o f   m u l t i l a y e r e d   s c a l e s   w h e r e   a n   i n n e r   c o m p a c t   o x i d e   g r o w s   a t   a   p a r a b o l i c   r a t e   a n d   i s   s i m u l t a n e o u s l y   o x i d i z e d   t o   a   h i g h e r   o x i d e   w i t h   n o   p r o t e c t i v e   p r o p e r t i e s .   I n   t h e   c a s e   o f   p l a n e   s y m m e t r y ,   t h e   g r o w t h   o f   t h e   t w o   l a y e r s   m a y   b e   d e s c r i b e d   b y   t h e   e q u a   t i o n s : 2 6 , 2 9   d y l d t   =   a l y     b   ( 1 5 )   d z l d t   =   b   ( 1 6 )   w h e r e   y   a n d   z   r e p r e s e n t   t h e   m a s s   g a i n   p e r   u n i t   a r e a   i n   t h e   c o m p a c t   l a y e r   a n d   i n   t h e   p o r o u s   l a y e r ;   a   a n d   b   a r e   c o n s t a n t s .   T h e   i n t e g r a t i o n   o f   t h e s e   e q u a t i o n s   g i v e s   t h e   g r o w t h   l a w   o f   e a c h   s c a l e :   y   =   a / b   l o g   ( l   b / a y ) ’     b t   ( 1 7 )   z   =   b t   ( 1 8 )   a n d   t h e   t o t a l   o x i d a t i o n   r a t e   o f   t h e   p a r a l i n e a r   o x i d a t i o n   i n   t e r m s   o f   m a s s   g a i n   x :   x   =   y   +   z   =   a / b   l o g   [   1     b / a   ( x     b t ) ] ’   ( 1 9 )   D u r i n g   i n i t i a l   o x i d a t i o n ,   t h e   p a r a b o l i c   g r o w t h   p r e   d o m i n a t e s   a n d   c o r r e s p o n d s   t o   t h e   g r o w t h   o f   y .   W h e n   a / y   =   b   t h e   g r o w t h   r a t e   o f   t h e   c o m p a c t   l a y e r   i s   e q u a l   t o   z e r o   a n d   y   r e a c h e s   a   m a x i m u m   v a l u e :   Y   m a x   =   a / b   ( 2 0 )   T h e n   t h e   r e a c t i o n   p r o c e e d s   a t   c o n s t a n t   r a t e :   d x l d t   =   s l y , , ,   =   b   ( 2 1 )   c o r r e s p o n d i n g   t o   l i n e a r   k i n e t i c s .   W h e n   t h i s   p s e u d o l i n e a r   r e g i m e   i s   e s t a b l i s h e d ,   t h e   f l o w   r a t e   o f   d i f f u s i n g   s p e c i e s   m a y   b e   d e r i v e d 3 ’   f r o m   t h e   w e l l k n o w n   F i c k ’ s   f i r s t   l a w :   d c u     =   D S   A C   d t     =   K , S   Y   m a x   ( 2 2 )   w h e r e   D ,   S   a n d   A C   a r e   r e s p e c t i v e l y   t h e   d i f f u s i o n   c o e f f i c i e n t ,   t h e   r e a c t i v e   a r e a ,   t h e   o x y g e n   c o n c e n   t r a t i o n   v a r i a t i o n   a n d   K .   a   r a t e   c o n s t a n t .   T h e r e f o r e ,   t h i s   e q u a t i o n   b e i n g   i d e n t i c a l   t o   ( 4 )   e x p l a i n s   w h y   l i n e a r   F ( a )   =   k t   p l o t s   h a v e   b e e n   o b s e r v e d .   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A n   i n t e r f a c i a l   l i m i t i n g   p r o c e s s   i s   s h o w n   t o   b e   d e v e l o p e d   d u r i n g   t h e   2 0   h o u r s   o x i d a t i o n   o f   t h e   T a C   g r a d e .   C o n c e r n i n g   T a , C ,   t h e   m e c h a n i s m   i s   c e r t a i n l y   m o r e   c o m p l e x .   T h e   p r e s e n c e   o f   a n   o x y c a r b i d e   l a y e r   T a C , O ,   h a s   b e e n   o b s e r v e d   a t   t h e   T a 2 C   T a 2 0 5   i n t e r f a c e   a n d   m a y   b e   e x p l a i n e d   b y   t h e   h e x a g o n a l   T a , C   s t r u c t u r e   w h i c h   p r e s e n t s   e m p t y   l a y e r s   o f   c a r b o n   a t o m s .   T h e s e   v a c a n c i e s   m a y   b e   o c c u p i e d   b y   t h e   o x y g e n   a t o m s   a n d   a   s t e a d y   s t a t e   l a y e r   t h i c k n e s s   i s   f o r m e d   i n d e p e n d e n t l y   o f   t i m e .   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S h i m a d a ,   S . ,   N i s h i s a k o ,   M .   a n d   I n a g a k i ,   M . ,   F o r m a t i o n   a n d   m i c r o s t r u c t u r e   o f   c a r b o n c o n t a i n i n g   o x i d e   s c a l e s   b y   o x i d a t i o n   o f   s i n g l e   c r y s t a l s   o f   z i r c o n i u m   c a r b i d e .   J .   S h i m a d a ,   S .   a n d   K o z e k i ,   M . ,   O x i d a t i o n   o f   T i C   a t   l o w   t e m p e r a t u r e s .   S h i m a d a ,   S . ,   A   k i n e t i c   a n d   t h e r m o a n a l y t i c a l   s t u d y   o n   o x i d a t i o n   o f   p o w d e r   a n d   s i n g l e c r y s t a l   s a m p l e s   f o r   n i o   b i u m   c a r b i d e .   K o f s t a d ,   P . ,   E l s e v i e r   S c i e n c e   P u b l i s h e r s ,   L o n d o n ,   1 9 8 8 .   L o r i e r s ,   J . ,   L o i   d ’ o x y d a t i o n   d u   d r i u m   m e t a l l i q u e .   G e n b a l i s a t i o n   a   d ’ a u t r e s   m e t a u x .   A c a d .   W e b b ,   W .   W . ,   N o r t o n ,   J .   T . ,   a n d   W a g n e r ,   C ,   O x i d a t i o n   o f   t u n g s t e n .   J .   1 1 .   S t r e i f f ,   R . ,   E t u d e   d e   l ’ o x y d a t i o n   d u   c a l c i u m   m a s s i f .   M t c a n i s m e s   d e   l a   r e a c t i o n .   B i l l y ,   M .   a n d   V a l e n s i ,   G . ,   R e a c t i o n   v e l o c i t y   o f   m e t a l   l o i d s   w i t h   m e t a l s ,   w h e n   t h e   i n t e r f a c i a l   r e a c t i o n s   a r e   a c c o m p a n i e d   b y   d i f f u s i o n :   a p p l i c a t i o n   t o   s u l f u r s i l v e r   r e a c t i o n .   J .   \\x0c']"
},{
  "_id": 32,
  "PDF": "Comparison of ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 oxidation mechanisms in air using LIF of BO2(g).pdf",
  "Text": "['Corrosion Science 163 (2020) 108278  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e lsev ie r .com / loca te /co rsc i  Comparison of ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 oxidation mechanisms in air using LIF of BO2(g)  T  V. Guérineaua, G. Vilmartb, N. Dorvalb, A. Julian-Jankowiaka,*  a DMAS, ONERA, Université Paris-Saclay, F-92322 Châtillon, France b DPHY, ONERA, Université Paris-Saclay, F-91123 Palaiseau, France  A R T I C L E  I N F O  Keywords: (A) ceramic (C) high temperature corrosion (C) oxidation (B) spectroscopy (B) laser induced fluorescence  1.  Introduction  A B S T R A C T  The oxidation behaviour of ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 is studied using the real-time Laser-Induced Fluorescence (LIF) detection of BO2(g) evaporated from samples heated up to 1873K in dry air. For each composition, the relevant steps of oxidation (silica formation, volatilization, etc.) can be precisely determined. Moreover, the influence of composition on the oxidation behaviour, and more precisely on the B2O3/SiO2 ratio in the glassy phase, can be understood and described by monitoring the LIF signal from the BO2(g) radical. Thus, this technique has confirmed the potential of the HfB2-SiC-Y2O3 composition as a Ultra High Temperature Ceramic (UHTC) material.  Ultra-High Temperature Ceramics (UHTC) are extensively investigated as materials for applications in very high temperatures (> 2073 K) and oxidising or corrosive atmospheres. Among this material family, UHTC borides, such as ZrB2 and HfB2-based ceramics, are particularly interesting since the melting points of these materials are high, and the mechanical, thermal and chemical properties are excellent. Hence, they might be ideal candidates for developing thermal protection systems on hypersonic aerospace vehicles or atmospheric reentry vehicles which require a good oxidation resistance under severe conditions. Thus, studying and understanding the oxidation resistance of such materials is essential. The influence of the SiC addition or of other additives (e.g. TaSi2, MoSi2) has been largely studied [1-5] so that, the oxidation behaviour of ZrB2-SiC and HfB2-SiC is relatively well known up to 2073 K in air [6-12]. A few studies have also been carried out at higher temperatures and under other atmospheres than air [13-19], for instance in the presence of humidity or in an oxyacetylene flame. Although the key role of the borosilicate layer [1,20,21] and finally of the ZrO2 or HfO2 layer [13,18,22] has been demonstrated, a very few studies attempted the in situ investigation of its behaviour during an oxidation test. The thermogravimetric analysis (TGA) [20] and the optical emission spectroscopy have been used elsewhere [23,24]. However, TGA only gives the sum of the weight gains and weight losses  and only provides insights about the chemical evolution of the borosilicate layer. Investigations by emission spectroscopy imply that the chemical species are formed in sufficient quantity in an excited state to be detected through the radiative deactivation process, thus making this technique slightly sensitive. Previously, we have proposed the Laser-Induced Fluorescence (LIF) technique to detect gaseous BO2(g) radicals above the sample during oxidation of ZrB2 and ZrB2-SiC samples [25]. These radical species come from the volatilization of B2O3 from the glassy layers. The LIF technique enables the detection of BO2 radicals with much better sensitivity, spatial and temporal resolutions than emission spectroscopy. Laser excitation of BO2 is performed at 547.3 nm to promote the A2Πu (0,0,0) X2Πg (0,0,0) vibronic band originating from the ground state. The subsequent fluorescence is detected via the A2Πu (0,0,0)→X2Πg (1,0,0) red-shifted band at 580 nm [25]. In contrast to the TGA analysis, oxidation detection is delayed since B2O3 has to be formed prior to the detection of BO2(g) [23]. However, the further evolutions of the glassy layer are markedly detected using the LIF technique such as the onset of the oxidation of SiC or the intense volatilization of the glassy layer at temperatures higher than 1500K. This paper presents a comparison of the oxidation mechanisms of ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 by using the same real-time LIF detection strategy of evaporated BO2 above the ZrB2, ZrB2-SiC samples heated in air, a technique which has been applied for the first time in our previous article [25]. The influence of the composition on the  ⁎ Corresponding author. E-mail addresses: vincent1guerineau@gmail.com (V. Guérineau), gautier.vilmart@onera.fr (G. Vilmart), nelly.dorval@onera.fr (N. Dorval), aurelie.jankowiak@onera.fr (A. Julian-Jankowiak).  https://doi.org/10.1016/j.corsci.2019.108278 Received 7 June 2019; Received in revised form 3 October 2019; Accepted 8 October 2019 Available online 15 October 2019 0010-938X/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'Table 1 Grade, purity and particle size of starting powders.  Powder  Provider  ZrB2 HfB2 SiC Y2O3  H.C. Starck H.C. Starck H.C. Starck Ampere Industrie  Grade  A A BF12 -  d50 (μm)  2.8 7.6 0.6 < 5  Purity  > 97.8 > 97 > 98.5 99.99  oxidation behaviour is described in details and the potential of the LIF technique for thermal oxidation detection is confirmed here in the case of the HfB2-based ceramics.  2. Materials and apparatus  2.1. Materials  Three UHTC compositions are studied: ZrB2 + 20 vol% SiC, labelled ZS, HfB2 + 20 vol% SiC, labelled HS and HfB2 + 20 vol% SiC + 3 vol% Y2O3, labelled HSY. Properties of the starting powders are listed in Table 1. After weighing, the powders are attrition-milled in cyclohexane using zirconia media (for ZS) or WC media (for HS and HSY), dried and sieved (50 μm). Fully dense samples (20-40 mm in diameter) are obtained using Spark Plasma Sintering (SPS, FCT HD25) in the MATEIS laboratory (Lyon, France). More details concerning manufacturing have already been reported elsewhere [13]. The bulk density and open porosity of these materials are measured by the Archimedes’ method. Then, the densification level is calculated as the ratio of the apparent density to the theoretical density of the powder mixture. Sintering conditions and densification levels of the samples are reported in Table 2.  2.2. LIF experiment  During the thermal oxidation runs in air of the samples heated with a continuous wave 2 kW CO2 laser, BO2(g) is detected in the evaporation plume by monitoring its fluorescence signal at 580 nm induced by laser excitation tuned to 547.3 nm as mentioned above [26]. A very detailed description of the set-up, the method and the experimental protocol were reported in our previous study [25]. The dedicated chamber is equipped with optical windows for laserbased measurements. The CO2 laser irradiates the target at normal incidence angle from the top of the chamber (Fig. 1a). The laser beam is focused on the target in order to irradiate the whole surface. The sample is mounted on an alumina holder in the chamber centre. The support can be translated upwards in order to adjust the height of the sample surface with respect to the probe laser beam axis. Surface temperature is measured with a bichromatic pyrometer in the 1273-2773 K range (Modline 6R-2565, Ircon). Synthetic air (Alphagaz 1) is injected in the chamber by means of a mass flow regulator, and is evacuated via a pumping system to maintain the pressure constant during the oxidation run. The probe laser source is a tunable dye laser (Quantel, TDL50) pumped by a pulsed Nd:YAG laser (Quantel, YG781, 10 Hz repletion rate). The pulse duration is 6 ns and the linewidth is 0.07 cm−1. The output energy is 15 mJ/pulse and is attenuated to 1 mJ using neutral  Table 2 : Sintering conditions and open porosity of the studied UHTC materials.  Composition  density filters. The probe laser beam is crossing the chamber horizontally through anti-reflection coated windows in the visible range. The beam is focused with a spherical lens (f = 600 mm) to obtain a 1 mm laser spot at the sample location in the chamber. During a thermal oxidation run, the laser energy is recorded at the chamber exit by a photodetector (PD10, Ophir) in order to monitor the beam attenuation through the evaporation plume (absorption pathway in Fig. 1b) Fluorescence light is collected at right angle from the probe laser beam axis (fluorescence pathway in Fig. 1b) by a 100 mm focal-length lens coupled to an optical fibre (1 mm in diameter). The fluorescence is imaged with a magnification of 1 onto a monochromator (Jobin-Yvon, H-20) equipped with a 1 mm wide exit slit (4 nm bandpass centred at 580 nm). The signal is then detected, amplified, time-integrated (34 ns temporal gate) and averaged (over 10 laser shots) by means of a photomultiplier tube (Photonis, XP2017B) and a Boxcar system (SRS, SR250) [25]. A Labview® interface is used to acquire integrated fluorescence signals via an acquisition PCI card (1.25 MHz, 16 bit) at 10 Hz.  2.3. Experimental protocol  Once the sample is positioned on the holder, the chamber is evacuated prior to synthetic air injection with a 1.3 l/min flow rate. The pressure is regulated to 0.1 MPa, then the heating of the sample and the recording of both the LIF signal and the laser transmission can start. The CO2-laser power is monitored manually, thus allowing to hold temperature steady to perform timely several adjustments and to ensure that the LIF signal is properly monitored [25]. The nature and the thickness of each oxidised layer were identified through ex situ examinations in the Scanning Electron Microscope (SEM, DSM 962) and by Energy Dispersive Spectroscopy (EDS) of both the surface and the polished cross-sections of the oxidised materials. In addition, TGA experiments (SETSYS Evolution 16/18, SETARAM) are carried out to measure the mass variations during oxidation of each composition. Samples (2 × 2 × 4 mm3 in size), are slowly heated (2 K/ min) from room temperature to 1823 K under a 50 ml/min air flow.  3. Results  The variation of amplitude of LIF and laser transmission signals as a function of time for ZS, HS and HSY samples is shown in Fig. 2. As for ZS samples [25], four key steps can be identified for HS and HSY samples and are numbered from ① to ④ in Fig. 2 but with noticeable differences in the amplitude of signal or in the temperature range associated with each step, as compared in Table 3. Step ① corresponds to the appearance of the LIF signal indicating that a detectable quantity of BO2(g) has been formed. This gaseous species comes from the volatilization of the B2O3(l) glassy layer formed at the surface of the sample (Eqs. (1) and (2)).  MeB2 + 5/2 O2 → MeO2(s) + B2O3 (l)  B2O3(l) → B2O3(g)  (1)  (2)  where Me = Zr or Hf. For all samples, weight gains are observed at lower temperature with TGA (1013 K for ZS and 948 K for HS and HSY) as shown in Fig. 3. Thus, as already observed, a threshold in boria amount and thus in B2O3(g) has to be overstepped before the appearance of the LIF signal. At this temperature (< 1473 K), SiC is not yet oxidised (Fig. 2) [3].  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  Fig. 1. (a) a schematic view of the chamber and (b) LIF set-up.  Thus, the following steady rise in the LIF signal between steps ① and ② is attributed to the enhanced volatilization of B2O3, leading to an increased quantity of BO2(g). This is confirmed by the slower increase in weight gain as measured by TGA (Fig. 3). An inversion of trend is observed at step ②: the LIF signal slowly decreases despite the continuous rise of sample heating, up to step ③, indicating that less BO2(g) is formed in this temperature range (Table 3). This striking observation is associated with the beginning of the passive oxidation of SiC through reaction 3. Therefore, the enrichment of the glassy phase with silica increases its stability (higher viscosity) compared to the previous boria glassy layer (Eq. (4)).  SiC + 4O2 → SiO2(s/l) + CO2(g)  SiO2 (l) + B2O3 (l) → {SiO2:B2O3}(l)  (3)  (4)  Several studies have already shown that a B2O3 concentration gradient is formed in the borosilicate glassy layer as preferential evaporation of B2O3 is observed at the surface [27-29]. This lower concentration, and thus the lower activity of B2O3 at the surface explains the LIF signal decrease [11,23,29]. Passive oxidation of SiC is also detected using TGA, but at a higher temperature. Between steps ③ and ④ small bumps in the LIF signal, immediately followed by a decrease in the signal are observed. This is called the “dynamic phase” and can be explained by a power peak which makes the oxide layer less stable. Thus, the B2O3-rich glassy layer, where the activity of B2O3 is higher, is exposed and, therefore, the volatilization of B2O3 is favoured. The glassy layer is thinner and less protective against oxidation, so that the oxidation of the material underneath is favoured and, as a consequence, quickly reforms a stable borosilicate glassy layer, and the LIF signal decreases. The subsequent signal bumps have the same origin (power peak) and the same consequences. Between each bump, the average intensity of the LIF signal increases as the temperature is continuously increasing. Compared to TGA, between 1300 and 1673 K, the weight gain increases strongly for HSY, and more smoothly for HS. ZS has a different TGA behaviour, especially at 1473 K. Instead of a decrease in the mass gain slope, which is observed for HS and HSY, there is an increase in this slope. This indicates a continuous oxidation of the materials. Finally, the high intensity LIF signal peak coupled with a dramatic drop in the laser transmission signal is observed in step ④. At that moment, the probe laser beam is strongly attenuated because of the formation of a dense smoke of oxide particles in the plume, and the LIF signal is no longer detected. This step is also triggered by a power peak, but at this high temperature, the volatilization of the glassy layer is very intense, and the active oxidation of SiC is also favoured via the Eq. (5).  SiC + O2(g) → SiO(g) + CO(g).  (5)  3  The drop in the transmission signal is linked to the catastrophic volatilization of the glassy layer leading to a high amount of B2O3 and SiO particles in the atmosphere just above the sample. However, at 1673 K, the mass gain measured by TGA increases noticeably for all the compositions (Fig. 3), which confirms that the volatilization of the borosilicate glassy phase promotes the oxidation of the material underneath. These general trends are observed for all compositions. However, some differences are noticeable, such as the temperatures (T1, T2 and T3) of the onset of the three steps ① to ③, the LIF signals between steps ③ and ④ and the LIF and transmission signals at step ④. These differences are discussed in the next section.  4. Discussion  4.1. Formation of  the boria glassy phase  Considering Table 3, it can be noticed that the temperature at which the LIF signal is first detected (T1) is slightly higher for ZS than for HS and HSY. Interestingly, weight gain is detected at lower temperatures in the TGA signal (around 1113 K for ZS and around 948 K for HS and HSY) indicating that the LIF signal BO2(g) is detected once a boria glassy layer is formed at the surface of the sample (Fig. 3). The later apparition of the B2O3 boria glassy phase and thus of the BO2(g) detection for ZS samples is observed using both experiments. This can be explained considering the volatility diagrams of ZrB2 and HfB2 [30,31], and the diagrams of Ellingham (Fig. 4). It is clearly shown that oxidation of HfB2 is favoured at a lower temperature than for ZrB2.  4.2. Formation of  the borosilicate layer  The temperature at which the LIF signal stops increasing (T2) is higher for ZS and HSY than for HS (˜100 K in difference). This phenomenon is worth noting since T2 marks the moment when SiO2(l) effectively integrates the boria network to form a borosilicate layer, making it more stable and less subject to volatilization. SEM micrographs of the three materials after oxidation are shown in Figs. 5 and 6. The oxidised samples all develop a classic 3-layer oxide scale above the bulk material, from top to bottom: a SiO2-rich glassy layer, a MeO2 layer (Me = Zr or Hf) and a SiC-depleted MeB2 layer. The SiC-depleted layer is formed by active oxidation of SiC (reaction 5), and the reaction product SiO(g) can diffuse towards the surface [32]. Once SiO(g) reaches the borosilicate layer, it can react with dissolved O2(g) to form SiO2(l), thus participating in the replenishment of the glassy layer [13,33,34] (Eq. (6)).  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  Fig. 2. Surface temperature (left axis), normalised LIF and transmission signals (right axis) vs. time of (a) ZrB2-SiC, (b) HfB2-SiC and (c) HfB2-SiC-Y2O3 materials during an oxidation test in dry air under 0.1 MPa. Key steps (numbered 1-4) are described in the text.  Table 3 Temperatures of the onset of the three key steps T1, T2 and T3. *extrapolated temperature.  Sample  Detection of LIF signal (T1, ①)  Step ② onset (T2).  Step ③ onset (T3)  ZS HS HSY  1263 K 1223 K 1223 K*  1553 K 1433 K 1523 K  1643 K 1573 K 1588 K  SiO(g) + ½ O2(g) → SiO2(l)  (6)  From the SEM observations, it should be noted that the SiC-depleted layer has a very different relative thickness depending on the nature of the sample. HS sample develops a very thick SiC-depleted HfB2 layer (˜500 μm, 90% of the total oxide scale), HSY develops a thick SiC-depleted HfB2 layer (˜250 μm, 60% of the total oxide scale) and ZS develops a comparatively thin SiC-depleted ZrB2 layer (˜60 μm, 25% of the total oxide scale) [10,29,32], as reported in Table 4.  4  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  Fig. 3. Mass variation vs. time of HS, ZS and HSY. Temperature ramp is 2 K /min. The TGA is performed under a 50 ml/min flux of air. Inset: a zoom on TGA signals between 800-1200 K.  Moreover, HS develops a very thin HfO2 layer compared to the SiCdepleted HfB2 layer or to other compositions. The composition of the glassy layer surface is directly dependent on the different reaction products that replenish it during the oxidation test. B2O3(l) only comes from the oxidation of HfB2 whereas SiO2(l) only comes from (directly or via active oxidation) the oxidation of SiC. Thus, the borosilicate layer developed by HS will likely be enriched in SiO2(l) compared to HSY and ZS. This enrichment leads to an early stabilisation of the glassy layer, and that is why step ② happens at a lower temperature for HS. For HSY, step ② happens at 1523 K, closer to the step ② of ZS (1553 K). First, HSY develops a thinner SiC-depleted layer compared to HS: this means that the borosilicate layer is comparatively enriched in B2O3, which delays the stabilisation effect provided by SiO2. Secondly, the glassy phase is able to solubilise Y2O3 during its convection to the surface [21,35,36,13]. Thus, the dissolution of Y2O3 in the borosilicate glass decreases its viscosity and, therefore, its capability to act as an oxygen diffusion barrier, since a low viscosity favours the diffusion of oxygen through the borosilicate layer. To form a SiC-depleted layer, the oxygen partial pressure PO2 should be low enough to oxidise SiC without oxidising HfB2 [32,10,37]. If the glassy layer is less protective towards oxygen diffusion, the formation of a SiC-depleted layer is hindered in favour of the oxidation of both SiC and HfB2. This explains why, when  Y2O3 is added to a HfB2-SiC composition, the stabilisation provided by SiO2(l) is triggered at a higher temperature. Then, comparing ZS and the two other compositions, ZS exhibits the thinnest SiC-depleted layer and the thickest glassy layer, indicating that the glassy layer formed at the top surface of ZS samples is B2O3-rich compared to the glassy phases developed by HS and HSY. Thus, the glassy layer of ZS is more permeable towards oxygen, thus promoting the oxidation of both SiC and ZrB2 (Fig. 5a and b). Another aspect that should be taken into consideration is that, at the experiment temperature (T < 1900 K), tetragonal-ZrO2 is a better oxygen conductor than monoclinic-HfO2 [38,39]. Thus, this point accentuates the fact that the active oxidation of SiC is not favoured in the ZS composition.  4.3. The dynamic phase  The beginning of step ③ always corresponds to a peak in the LIF signal (Fig. 2) and a strong increase in the weight gain (Fig. 3). In this temperature range (> 1573 K) the volatilization of B2O3 becomes very intense [20] and explains the constant increase in the average level of LIF signal observed in Fig. 2. However, in the present cases, the volatilization of boria is attenuated by the presence of silica in the glassy layer and the composition of the sample induces some differences in the  Fig. 4. Diagram of Ellingham for Si, Zr and Hf elements.  5  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  Fig. 5. SEM micrographs of polished cross-sections of oxidised samples: (a and b) ZS, (c and d) HS and (e and f) HSY. The different layers are (1) borosilicate glass, (2) MeO2 (Me = Zr or Hf), (3) SiC-depleted MeB2 and (4) bulk material.  Fig. 6. SEM micrographs of the surfaces of oxidised samples: (a and b) ZS, (c and d) HS and (e and f) HSY materials. In grey, the borosilicate layer and in white, zirconia or hafnia inclusions and dendritic structures. (f) Recrystallised HfO2 is omnipresent (in white) as dendrites and crystallites in the borosilicate layer (grey).  6  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  Table 4 Thickness of each layer in the oxidised layers for all (average values from 10 measurements).  the tested compositions  peak, the borosilicate layer is even thinner, and this triggers the oxidation of the underlying HfB2, which will refill the borosilicate glass at the surface.  SiC-depleted layer (μm)  Glassy phase layer (μm)  Total oxidised layer (μm)  4.4.  Intense volatilization of B2O3  ZS HS HSY  60 500 250  180 55 170  240 555 420  LIF and transmission signals. For HSY, once the dynamic phase has begun, the average level of the LIF signal remains almost constant, despite the continuous increase in temperature. In parallel, the transmission slowly decreases from 1473 to 1823 K. This means that there is a continuous formation of B2O3 droplets and, therefore, a steady volatilization of the borosilicate phase. But, the HSY material developed a very thick borosilicate layer containing a large amount of recrystallised HfO2 particles (Fig. 6b). The thickness of the glassy layer explains its resilience, i.e. its capacity to recover an average LIF signal level after the power peak. Moreover, Talmy et al. [2] have shown that the presence of immiscible phases in a glassy layer could increase its viscosity and, therefore, limit the diffusion of oxygen. During its convection, the B2O3-rich glass is able to solubilise Y2O3 and HfO2. When reaching the surface, B2O3 preferentially evaporates, so that Y2O3 and HfO2 attain their saturation concentration and precipitate. Therefore, the continuous volatilization of B2O3 throughout the oxidation test favours the continuous precipitation of the solubilised components, which, in turn, increases the stability of the glassy layer. For ZS, during the dynamic phase, the average level of the LIF signal increases after each peak (Fig. 2). For this sample, the stabilisation provided by SiO2 is less effective since the glassy layer is B2O3-rich, and the volatilization of B2O3 is, therefore, thermo-activated throughout the dynamic phase. The laser transmission signal also decreases steadily; consequently, there is a continuous formation of B2O3 droplets. Compared to HSY, the LIF signal takes more time to reach the low level after each signal peak (˜180 s compared to ˜60 s for HSY). This means that a longer time is required for the sample to regenerate a glassy layer with a good stability. This is mostly due to the fact that the glassy layer of the oxidised ZS sample is B2O3-rich, which means that it is more subjected to volatilization, and that a larger quantity is required in order to get the same protection, due to the lower viscosity of B2O3-rich borosilicate glass. On the contrary, the glassy layer developed by HSY contains relatively more SiO2, so that it is more stable, and the material developed a thick SiC-depleted layer, which means that the replenishment of the layer is made with a SiO2-richer glass, explaining why the glassy layer more rapidly attains stability. The LIF signal behaviour of HS during oxidation is similar to that of ZS. Each LIF signal peak is followed by an increase in the average LIF signal level. But, contrary to the other samples, the laser transmission signal was almost constant throughout the experiment. The LIF peaks are also very sharp. To explain these observations, it should be remembered that the borosilicate glass formed by HS is likely very SiO2rich, since the SiC-depleted zone is 9 times thicker than the HfO2 layer. The deficiency in B2O3 explains that there is no or not enough formation of B2O3 droplets and, therefore, the transmission level is unchanged. It can be noticed that the LIF signal peak at 2100s (1783 K) is accompanied by a short drop in the transmission level, thus meaning that, at this temperature, some B2O3 droplets are formed in a short time. As evidenced in Fig. 5, the borosilicate layer is likely thin and SiO2-rich during the test. Contrary to the HSY sample, where the height of LIF signal peaks was about a half of the average LIF signal level, the LIF signal peaks observed for HS are much more prominent, and can be ten times higher than the average LIF signal level. Due to the B2O3 deficiency and, therefore, the thinness of the HS borosilicate layer, this glassy phase is much more sensitive to a power peak. After the power  Step ④ also illustrated some differences in the oxidation behaviour of the different materials. For the ZS sample, the final LIF signal peak was accompanied by a drastic drop in transmission. For the HS material, this step is characterised by a moderate LIF signal peak and a decrease in the transmission signal. And, finally, for HSY, there is a high intensity LIF signal peak, but there is only a small laser transmission decrease (from 0.7 to 0.6). Regarding ZS sample, the LIF and transmission signals are associated with a catastrophic volatilization of the B2O3-rich borosilicate layer. This is confirmed by the final microstructure of the oxidised ZS sample (Fig. 5a), where the glassy layer is very thin or non-existent. However, SEM micrographs of the surface of the oxidised ZS sample (Fig. 6a and b) clearly show two features: the first is that, despite the thinness of the glassy layer, it still almost covers the whole surface of the sample, and the second is that the glassy layer is filled with recrystallized ZrO2 (small particles and large dendritic structures). The crystallites in the glassy layer constitute an indirect proof of the intense volatilization of B2O3. The fact that the borosilicate layer covers the entire surface reveals the low viscosity of this layer, probably due to its poverty in SiO2. The step ④ of HSY is also characterised by an intense LIF signal peak, but not a drastic transmission drop. As explained earlier, the thickness of the glassy layer developed by HSY allows the borosilicate glass to be more resilient towards the power peak. Even if the LIF signal of BO2 increases, there are not enough B2O3 droplets formed to drastically lower the laser transmission and, therefore, the volatilization cannot be qualified as catastrophic. This is confirmed by the SEM observations; the glassy layer is thick (Fig. 5e and f) and covers the whole surface (Fig. 6e and f). As in the case of ZS, the overwhelming presence of crystallites and dendritic structures of Y2O3-doped HfO2 at the surface constitutes an evidence of the constant volatilization of B2O3 during the oxidation test. The case of HS is very different. As explained above, the step ④ of HS is not characterised by an intense LIF signal peak. But a closer examination of the transmission signal reveals that there is a very short but intense drop at 2400s (1833 K). SEM observations of the HS microstructure showed that the glassy layer was thin (Fig. 5c and d) and did not cover the whole surface of the sample (Fig. 6c and d). The fact that there is simply not enough B2O3 in the remaining borosilicate layer to generate enough detectable BO2 can explain that there is no intense LIF signal at the end. The lack of B2O3 means that the viscosity of the borosilicate layer increases, thus, explaining that the glassy layer does not cover the entire surface of the sample.  5. Conclusion  The oxidation behaviours of ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 materials were studied using the real-time LIF detection of evaporated BO2(g), above the samples heated up to 1873 K in a dry air atmosphere. For each composition, the starting point of the influence of SiC on the stability of the glassy layer and the catastrophic volatilization of the glassy layer can be precisely determined. Moreover, the influence of the composition on the oxidation behaviour can be understood and described using this technique. In order to form a stable glassy layer, the B2O3/SiO2 ratio must be finely tuned. If this ratio is too high, as it seems to be the case for ZS, then the glassy layer is not stabilised enough by SiO2, and the glassy layer will completely volatilize at ˜1873 K. If it is too low, for instance for the HS material, then there is not enough \"matter\" to create a thick glassy layer that will completely cover the surface of the sample. On the contrary, a good compromise is achieved for HSY: this sample is not  7  \\x0c', 'V. Guérineau, et al.  Corrosion Science 163 (2020) 108278  [17]  [14]  [12]  Am. Ceram. Soc. 98 (2015) 1673-1683. [11] K. Shugart, W. Jennings, E. 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Dorval, In situ study of the oxidation of ZrB2 and ZrB2-SiC materials by monitoring the LIF signal of BO2 radicals, Corros. Sci. 148 (2019) 31-38. J.W.C. Johns, The absorption spectrum of BO2, Can. J. Phys. 39 (1961) 1738-1768. S.N. Karlsdottir, J.W. Halloran, Oxidation of ZrB2-SiC: influence of SiC content on solid and liquid oxide phase formation, J. Am. Ceram. Soc. 92 (2009) 481-486. S.N. Karlsdottir, J.W. Halloran, C.E. Henderson, Convection patterns in liquid oxide films on ZrB2-SiC composites oxidized at a high temperature, J. Am. Ceram. Soc. 90 (2007) 2863-2867. S.N. Karlsdottir, J.W. Halloran, Formation of oxide scales on zirconium diboride-silicon carbide composites during oxidation: relation of subscale recession to liquid oxide flow, J. Am. Ceram. Soc. 91 (2008) 3652-3658. [30] W.G. Fahrenholtz, The ZrB2 volatility diagram, J. Am. Ceram. Soc. 88 (2005) 3509-3512. [31] H. Li, L. Zhang, Q. Zeng, L. Cheng, Thermodynamic calculation of HfB2 volatility diagram, J. Phase Equilibria Diffus. 32 (2011) 422-427. [32] W.G. Fahrenholtz, Thermodynamic analysis of ZrB2-SiC oxidation: SiC-depleted region, J. Am. Ceram. Soc. 90 (2007) 143-148. T.A. Parthasarathy, R.A. Rapp, M. Opeka, M.K. Cinibulk, Modeling oxidation kinetics of SiC-containing refractory diborides, J. Am. Ceram. Soc. 95 (2012) 338-349. E.P. Simonenko, D.V. Sevast’yanov, N.P. Simonenko, V.G. Sevast’yanov, N.T. Kuznetsov, Promising ultra-high-temperature ceramic materials for aerospace applications, Russ. J. Inorg. Chem. 58 (2013) 1669-1693. [35] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. Z. Tian, Y. Zhu, W. Huang, Effect of yttrium oxide on the viscosity and crystallization of borosilicate glass at high temperature, Adv. Mater. Res. 750-752 (2013) 974-977. E.J. Opila, J.L. Smialek, R.C. Robinson, D.S. Fox, N.S. Jacobson, SiC recession caused by SiO2 scale volatility under combustion conditions: II, thermodynamics and gaseous-diffusion model, J. Am. Ceram. Soc. 82 (1999) 1826-1834. [38] M.F. Trubelja, V.S. Stubican, Ionic conductivity of the fluorite-type hafnia-RE2O3 solid solutions, J. Am. Ceram. Soc. 74 (1991) 2489-2494. [39] G.S. Corman, V.S. Stubican, Phase equilibria and ionic conductivity in the system ZrO2-Yb2O3-Y2O3, J. Am. Ceram. Soc. 68 (1985) 174-181.  formation of a  [26] [27]  [34]  [28]  [29]  [33]  [36]  [37]  only able to generate a thick borosilicate layer, but the preferential volatilization of B2O3 will trigger the precipitation of HfO2 particles, that will in turn increase the stability of the glassy layer in the considered temperature range. As it was already shown, LIF is a promising method, which can be quantitative, to study in situ, and thus continuously, the oxidation behaviour of UHTC. Moreover, the influence of the composition and its ability to form a protective layer can be studied and determined. Thus, the interest of the LIF technique in the detection of thermal oxidation is clearly confirmed.  Authors’ contributions  V. Guérineau is the PhD student who has manufactured the samples, tested them, performed the post-test analyses and written the article. N. Dorval and G. Vilmart are the specialists of the LIF technique and have helped to modify the bench and prepare the experiments. V. Guérineau, N. Dorval and G. Vilmart have performed all the LIF experiments. A. Julian-Jankowiak is the PhD supervisor, she has helped to modify the bench, to analyse the results and to write the paper. All the authors have read and approved the final version of article.  this  Data availability  The raw/processed data required to reproduce these findings cannot be shared at this time due to legal or ethical reasons.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgements  The authors would like to thank M. Bejet and T. Schmid (ONERA) for modifying the oxidation chamber and, finally, Dr P. Beauchêne and Dr. R. Valle (ONERA) for their kind and efficient assistance.  References  [3]  [2]  [1]  [4]  [5]  E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2and HfB2-based ultra-high temperature ceramics: effect of Ta additions, J. Mater. Sci. 39 (2004) 5969-5977. I.G. Talmy, J. Zaykoski, M.M. Opeka, S. Dallek, Oxidation of ZrB2 ceramics modified with SiC and group IV-VI transition metal diborides, Electrochem. Soc. Proc. 12 (2001) 144-158. P.A. Williams, R. Sakidja, J.H. Perepezko, P. Ritt, Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content, J. Eur. Ceram. Soc. 14 (2012) 3875-3883. P. Zhang, P. Hu, X. Zhang, J. Han, S. Meng, Processing and characterization of ZrB2SiCW ultra-high temperature ceramics, J. Alloys. Compd. 472 (2009) 358-362. S.C. Zhang, W.G. Fahrenholtz, G.E. Hilmas, Oxidation of ZrB2 and ZrB2-SiC ceramics with tungsten additions. High temperature corrosion and materials chemistry, Electrochem. Soc. Trans. 16 (2008) 137-145. [6] C.M. Carney, P. Mogilvesky, T.A. Parthasarathy, Oxidation behavior of zirconium diboride silicon carbide produced by the spark plasma sintering method, J. Am. Ceram. Soc. 92 (2009) 2046-2052. [7] W.G. Fahrenholtz, G.E. Hilmas, Oxidation of ultra-high temperature transition metal diboride ceramics, Int. Mater. Rev. 57 (2012) 61-72. F. Monteverde, The thermal stability in air of hot-pressed diboride matrix composites for uses at ultra-high temperatures, Corros. Sci. 47 (2005) 2020-2033. [9] C.M. Carney, Oxidation resistance of hafnium diboride-silicon carbide from 1400 to 2000 °C, J. Mater. Sci. 44 (2009) 5673-5681. [10] K. Shugart, E. Opila, SiC depletion in ZrB2-30 vol% SiC at ultrahigh temperatures, J.  [8]  8  \\x0c']"
},{
  "_id": 33,
  "PDF": "COMPRESSIVE STRENGTH DEGRADATION IN ZrB2-SiC AND ZrB2-SiC-C ULTRA HIGH TEMPERATURE COMPOSITES.pdf",
  "Text": "['Mechanical Properties and Performance of  Engineering Ceramics and Composites IV   Edited by Dileep Singh and Waltraud M. Kriven  Copyright 0 2010 The American Ceramic Society   COMPRESSIVE STRENGTH DEGRADATION IN ZrB2-SiC AND ZrB2-SiC-C ULTRA HIGH  TEMPERATURE COMPOSITES   J. Ramirez-Rico, M. A. Bautista, J. Martinez-Fernandez  Dpto. Fisica de la Materia Condensada-ICMSE  Universidad de Sevilla-CSIC  Avda. Reina Mercedes, s/n, 41012 Sevilla, Spain   M. Singh  Ohio Aerospace Institute  MS 106-5, Ceramics Branch  NASA Glenn Research Center  Cleveland, OH 44135-3191   ABSTRACT  The high melting point of refractory metal diborides makes them promising materials for  ultra high temperature applications. In this work, we study the compressive strength of two  ZrB2-SiC and ZrB2-SiC-C composites. Samples have been studied  in compression at room  temperature, 1400°C and 1550°C, in atmospheric air.  The degradation of the mechanical  properties as a result of atmospheric air exposure at high temperatures has also been studied as a  function of exposure time. The presence of C is detrimental to the compressive strength, as  carbon burns out at high temperatures in air. After exposure to air at high temperatures an  external S i02 layer is formed, below which ZrB2 oxidizes to Z r02. A reduction of 30% in room  temperature strength occurs after 16-24h of exposure to air at 1400°C for the ZrB2-SiC material,  while for  the ZrB2-SiC-C composition  this reduction  is observed after  less  than 6h. The  thickness of the oxide layer has been measured and the oxidation process is discussed in terms of  the existing models.   INTRODUCTION  The high melting point of refractory metal diborides coupled with their ability to form  refractory oxide scales give these materials the capacity to withstand temperatures in the 19002500°C range. These Ultra-High Temperature Ceramics (UHTC) were developed in the 1960s1.  Fenter2 provides a comprehensive review of the work accomplished  in the 1960s and early  1970s. Additions of Silicon Carbide are used to enhance oxidation resistance and limit diboride  grain growth3\"6. Carbon  is also sometimes used as an additive  to enhance  thermal stress  resistance7\\'8.  These materials offer a good combination of properties that make them candidates for  airframe  leading edges on sharp bodied reentry vehicles9. UHTCs have some potential  to  perform well in such applications\\' environment, i.e. air at low pressure. However, for hypersonic  flight in the upper atmosphere one must recognize that stagnation pressures can be greater than  one atmosphere. Some interest has also been shown in these materials for single use propulsion  applications10.  Major  improvements  in the manufacturing and characterization of ZrB2 .materials and  composites have been put forward in recent years, and now several important aspects of their  properties and processing are well understood4\\' 1M 9. However, the study of high temperature   127   \\x0c', \"Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   properties has been mostly limited to oxidation behavior, area which is also well understood '  6. To the best of our knowledge, very few studies of high-temperature mechanical properties  exist7·10·27.  In this work, we study the mechanical strength of ZrB2-SiC and ZrB2-SiC-C composites.  Samples were studied in compression at room temperature, 1400°C and 1550°C, in atmospheric  air. The degradation of the mechanical properties as a result of atmospheric air exposure at high  temperatures has also been studied as a function of exposure time. The microstructure and  composition of the as-fabricated and tested materials has been studied by means of Scanning  Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS).   MATERIALS AND METHODS  Samples of ZrB2-SiC composites were fabricated by uniaxial hot-pressing by Materials  and Machines, Inc., Tucson, AZ, using a procedure previously described . Two polycrystalline  composites were studied, containing variable amounts of ZrB2 and SiC; one of them also  contained C. In both cases the ZrB2/SiC volume ratio was 4. The nominal compositions of the  materials investigated are summarized in Table 1. ZrB2 (dso= 3-5 μπι) and α-SiC (dso= 1.4 μπι)  powders were obtained from H. C. Starck, and C powders were obtained from Asbury Graphite  Mills. Densities obtained were 5.57 g/cm3 (99.9% theoretical density) for ZS and 4.50 g/cm3  (99.0% theoretical density).  Compression tests were carried out on an electromechanical universal testing machine  with a furnace attached to its frame, at constant cross-head displacement rate. Load was applied  using alumina rods with SiC pads. Samples were cut into parallelepiped shape using a low speed  diamond saw. Nominal sample dimensions were 3x3x5 mm, and the load was applied to the  longest dimension. High temperature mechanical tests were conducted at room temperature,  1400°C and 1550°C. Several samples were exposed to oxidation by annealing at 1400°C in  atmospheric air  in a tube furnace, and exposure  times ranged from 6h to 48h. The room  temperature strength was measured after oxidation, to study the degradation of the mechanical  properties after exposure to an oxidizing environment. At least three samples were studied at  each temperature or exposure time at 1400 °C. Error bars throughout represent one standard  deviation.  Microstructural studies were carried out using SEM and EDS techniques, in both the asfabricated and  tested specimens. Samples were prepared using conventional metallographic  techniques which involved cutting, grinding and lapping. A conductive coating of either carbon  or gold was applied to the specimens prior to observation.   Table 1. Nominal composition and designation of the two composites studied in this work   Acronym  ZS   zsc   ZrB2 (vol. %)  80  56   SiC (vol. %)  20  14   C (vol. %)  0  30   RESULTS AND DISCUSSION   Microstructure of as-received specimens   1 28   ·   Mechanical Properties and Performance of Engineering Ceramics and Composites IV   \\x0c\", 'Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   Figure 1 shows the as fabricated microstructure of the two compositions studied. Our  observations match those reported in Ref. 7. The ZS composite appears to be fully dense, while  the ZSC suffered from significant grain pullout during polishing. This is attributed to the weak C  bonding to the ZrB2 and SiC, phases, which results in removal of the C phase during polishing.  In ZS, ZrB2 grains are equiaxed with reported grain size in the 6-12 μπι range, while the SiC  grains are elongated with sizes of approximately 1.5-3 μηι thick by 3-11 μιη wide/long7. In ZSC,  the grain pullout during polishing made the estimation of grain size difficult, although it can be  seen from Figure 1 that ZrB2 grain size is smaller in ZSC than in ZS. It should be noted that, at  least for the ZS composite, the grain size is close to the critical grain size for microcracking due  to the anisotropie thermal expansion coefficient of ZrB2, which has been reported to be around  15 μηι28.   Figure 1. Microstructure of the as-received materials. (ZS) ZrB2 20% vol. SiC. (ZSC) ZrB2 30% vol. C 14% vol. SiC. In the ZSC material, grain pullout and C removal during the  metallographic preparation is evident.   High Temperature Strength  The high temperature compressive strength was measured at 1400°C and 1550°C in  atmospheric air for both the ZS and ZSC compositions. Figure 2 shows representative stressdisplacement curves for the compositions and temperatures studied, while Figure 3 shows the  average compressive strength as a function of temperature. The ZS composite shows higher  compressive strength than the carbon containing ZSC. This is attributed to both the weak carbon  bonding to the other phases, evidenced as grain pullout in the SEM observations, and the burnout  of carbon at high temperature in air. This creates porosity and also produces channels through  which air can enter, oxidizing the ZrB2 phase not only on the surface but also inside the sample.  At 1550°C, creep of the SiC pads used as a protection to the alumina rods could be observed at  high stresses for the ZS material. For the ZSC material this effect was not observed due the much  lower maximum applied stress.  The relatively low values for the compressive strength in ZS at room temperature could  be attributed to the large grain size, which probably induces some microcracking in the ZrB2  matrix. This can also explain why at room temperature the composite containing carbon is  stronger, as it has a smaller grain size due to the presence of grain-growth inhibiting carbon.  Additionally,  the weak bonding of C to ZrB2 evidenced as grain pullout  in Figure 1 can  contribute to the relaxation of microcrack-inducing residual stresses developed upon cooling. At  high temperatures however, carbon burn-out is responsible for the lower strength observed in  ZSC when compared to ZS.   Mechanical Properties and Performance of Engineering Ceramics and Composites IV   ·   129   \\x0c', 'Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   0.00   0.05   0.10   Strain   0.15   0.20   Figure 2. Typical Stress-Displacement curves for ZS and ZSC at room temperature, 1400 °C and  1550°C.   0   200 400 600 800 1000 1200 1400 1600   Temperature (5Q   Figure 3. Compressive strengths of ZS and ZSC as a function of temperature.   Room Temperature Strength and Degradation  The degradation in the mechanical properties was studied as a function of exposure time to  atmospheric air at 1400°C. Figure 4 shows some selected stress-strain curves for ZS (left) and  ZSC (right), for exposure times ranging from 0 to 48h at 1400°C in atmospheric air. These  results are summarized  in figure 5, where the room temperature strength of the materials is  plotted as a function of exposure time. For ZS, a reduction of 50% in strength occurred between  16h and 24h, while for ZSC less than 6 hours were needed. This can be explained considering  that C burns in the oxidizing atmosphere, creating pores and channels through which air can  permeate.  It can be seen in Figure 5 vthat the strength of ZS improves after lh exposure to atmospheric  air at 1400 °C. This effect can probably be explained as a blunting of defects and anisotropyinduced microcracks by the silica glass layer that is formed during oxidation, as will be described  in the next section.   130   ·   Mechanical Properties and Performance of Engineering Ceramics and Composites IV   \\x0c', 'Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   0.00   0.02   0.04   0.06   0.08   0.10   0.12   0.00   0.02   0.04   0.06   0.08   0.10   0.1   Cross-head ctisplaoement (m r r /mm)   Cross-head displacement (mrn/rnm)   Figure 4. Typical stress-displacement behavior for ZS (left) and ZSC (right) samples oxidized in  atmospheric air at 1400 °C for different holding times.   0   10   20   30   40   50   Exposure Time at 140CPC in Air (hours)   Figure 5. Compressive strength degradation after oxidation in atmospheric air at 1400 °C. For ZS  samples, 50% strength is lost after 16-24h of exposure time, while for ZSC samples 50% of  the  strength is lost during the first 0-6h of oxidation.   Oxide layer formation  The microstructure and thickness of the oxide layers formed after exposure to atmospheric  air at 1400°C was studied for both ZS and ZSC samples, which were cut and polished for  observation in the SEM. Figure 6 shows a micrograph and compositional maps for a ZS sample  annealed for 1 hour. The outer, Si and O rich layer can be concluded to be S1O2, while the  intermediate layer, which is O rich and B poor, is ZrC>2. Similar conclusions can be drawn from  Figure 7, which represents compositional maps for a ZSC sample annealed for 24h. Again, the  outer layer is mainly composed of S1O2, and an intermediate layer of Zr(>2 separates the former  layer and the bulk interior of the sample.   Mechan ical Properties and Per formance of Engineering Ce ram ics and Compos ites IV   ·   1 31   \\x0c', 'Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   Figure 6. Typical microstructure of ZS samples after oxidation for lh at 1400 °C. (Top) BSE  micrograph of a section. (Bottom) O, Si, Zr and B EDS maps. An outer, glassy S1O2 layer is  followed by a porous ZrC>2 layer.   Figure 7. Characteristic EDS composition maps of ZSC after oxidation for 24h at 1400 °C.   These observations were confirmed by quantitative analysis, of which an example is  presented. Figure 8 is a cross section of a ZSC sample annealed for 24h at 1400°C in air, and  contains several of the features previously observed. The microstructure of the oxide layer can be  divided into four different zones or regions. The first, outer one is composed of Z r02 grains  embedded in a glassy S1O2 matrix, as can be deduced from the elemental composition. The zone  labeled as number 2 is composed only of S i02, while the zone labeled as 3 is composed almost  exclusively of ZrC>2 and contains no Si. The fourth zone corresponds to the composition of the  bulk, as-fabricated material. For ease of comparison, raw spectra obtained from all four different  zones are depicted.   132   ■ Mechanical Properties and Performance of Engineering Ceramics and Composites IV   \\x0c', \"Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   <:   Figure 8. (Left) Characteristic microstructure of ZSC samples after 24h exposure to oxidizing  atmosphere at 1400°C. Several zones can be distinguished. (Right) Relevant EDS spectra  obtained from the four zones depicted.   Biergy(k^/)   5   10   15   20   25   Hddngtime(h)   5   10   15   '   20   Holding t ime (h)   Figure 9. Reaction layer thickness as a function of holding time for ZSC (left) and ZS (right)  exposed to atmospheric air at 1400 °C.   These observations confirm the oxidation process already outlined in previous references,  such as Re f s .7'8'2 3'2 4'2 6. Both ZrB2 and SiC are oxidized, producing B2O3 that evaporates at high  temperatures. The S1O2 formed, which is liquid at the studied temperatures, is expelled towards  the surface of the sample by capillary forces, and acts as a protective layer. The intermediate  layer is thus composed mostly of Z r02 and pores that allow for oxygen permeation. It is thus  expected that the ZSC samples, containing carbon  that burns out at high temperature, will  oxidize at a faster rate because of the porosity produced during carbon combustion.  To ascertain  this effect,  the  thickness of both  the S1O2 and Z1O2 oxide  layers were  measured as a function of annealing time, for both ZS and ZSC samples. These results are  presented in Figure 9. It can be seen that the reaction rate is faster for samples containing C, and  that the thickness of the oxide layers is smaller for ZS samples.   Mechanical Properties and Performance of Engineering Ceramics and Composites IV   ·   133   \\x0c\", 'Compressive Strength Degradation in ZrB2-SiC and ZrB2-SiC-C Composites   CONCLUS IONS  It has been shown  that the addition of C to a ZrB2-SiC compos i te  is detrimental  to the  high  tempe ra tu re mechan ical  p rope r t ies,  since  the  degrada t ion  in  strength  and  oxidation  res is tance observed coun ters any poss ib le  imp rovement  to  the  thermal shock  resistance. The  observed s treng ths are low when compa red to o ther ma ter ia ls in the literature, espec ia l ly for finer  grained ma ter ia ls, wh ich can be attributed to the grain size of Z rB2 be ing c lose to the critical size  for m ic roc rack ing.  The ZS compos i t ion shows the best performance of the two compos i t ions  studied, both  in  terms of strength and oxidation  res is tance. The ZS material can withstand  exposures up to 24h  in air at 1400°C before  its comp ress ive strength  falls be low 50% when  compared  to the as-fabricated ma te r ia l, but the ZSC compos i t ion  is probab ly not suitable  for  application if no pro tec t ive coa t ing is app l ied.   ACKNOWLEDGEMENTS  This material  is based upon work  suppor ted by  the European Office  of Aerospace  Research  and Deve lopmen t, A ir Fo rce Office  of Scientific Research, Air Force Research  Labora tory, under Grant No. FA8655 -07 -1 -3087. 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},{
  "_id": 34,
  "PDF": "Copy of Separating Test Artifacts from Material Behavior in the Oxidation Studies of HfB2-SiC at 2000 degrees C and Above.pdf",
  "Text": "['Int. J. Appl. Ceram. Technol., 10 [2] 293-300 (2013)  DOI:10.1111/j.1744-7402.2011.02730.x  Separating Test Artifacts from Material Behavior in the -SiC at 2000°C and Above Oxidation Studies of HfB2  Carmen M. Carney* and Triplicane A. Parthasarathy  UES, Inc, Dayton, Ohio 45432  Michael K. Cinibulk  Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright Patterson Air Force  Base, Ohio 45433  Oxidation characteristics of HfB2-15 vol% SiC prepared by field-assisted sintering was examined at 2000°C by heating  it  in a zirconia-resistance furnace and by direct electrical  resistance heating of  the sample. Limitations of  the material and the  direct electrical resistance heating apparatus were explored by heating samples multiple times and to temperatures 2300°C. Oxide scales that developed at 2000°C from both methods were similar  in that  they consisted of a SiO2/HfO2 outer  in excess of  layer, a porous HfO2 layer, and a HfB2 layer depleted of SiC. But  they differed in scale thicknesses,  impurities present,  scale  morphology/complexity. Possible test artifacts are discussed.  Introduction  resistance compared with the diborides alone near 1600° C.1-3 At temperatures below 1800-2000°C, a refractory  Characterization of  the oxide scale formed on ultra  porous metal oxide scale is formed that  is protected by a  high temperature ceramics (UHTCs) has been a topic of  glassy silica scale. However, as temperatures are increased  intense study over the past decade. In particular, compos the protective  silica becomes  less  viscous  and thus  less  ite  systems  of  diborides  of  hafnium or  zirconium  protective. Hypersonic flight will  require  leading edges  with SiC have been studied for their improved oxidation  and nose  cone  components  and cooling to temperatures  to withstand rapid heating in excess of 2000°C under  This  study was  supported by the United States Air Force Contract # FA8650-10-D-5226.  *carmen.carney@wpafb.af.mil © 2013 The American Ceramic Society  shear stresses imparted by air flow. In addition, the envi ronment within the boundary layer near the component  \\x0c', 'will be comprised of a fraction of dissociated atoms depending on velocity, altitude, and other factors.4,5 Dis sociated oxygen can alter  the boundary between active  and passive oxidation of SiC and thus kinetics.6 These  influence oxida tion  temperatures  and  conditions  are  unattainable  in traditional molybdenum disilicide  ele ment  furnaces whereas  conventionally  accepted arc  jet  testing is expensive and the primary oxidant is dissociated  oxygen, so new methods of testing have been developed. Methods such as oxyacetylene torch heating,7-9 laser heatthe sample itself,11-13 ing,10 direct resistance heating of and scramjet simulators14 are being developed. The rapid  heating profiles and higher  temperatures attainable with  these tests may lead to different oxide morphologies and  performance  than  those  observed with  furnace-heated  samples. It is imperative that a correlation between differ ent  testing methods is made so that samples prepared by  different  exposure methods may be compared. To this -SiC samples were heated in air at 2000°C  end, HfB2  using a  zirconia-resistance  furnace  and direct  resistance  heating of  the sample and the resulting oxidation prod ucts were compared. In addition,  the limits of resistance  heating including multiple cycles and maximum temper atures were examined.  Experimental Procedure  Commercially available HfB2 (Materion, Milwaukee, WI, 99.9%, 45 lm) and b-SiC (Materion, 99.9%, 1 lm) were used to prepare HfB2-15 vol% SiC (HS).  The powder mixtures were ball milled in isopropanol  for 24 h with SiC grinding media, dried at  room tem perature, and subsequently dry milled for 12 h. Typical  weight  loss of  the SiC grinding media after milling was  0.2 mg (0.2 wt% of the total batch). The were sieved through an 80-mesh (177 lm)  powders  screen.  A quantity  of  150 g  of  the milled  powders was  loaded  into  a  60-mm diam.  graphite  die. A layer  of  BN and graphite  foil  separated the  powder  from the  die with the powder  in contact with the graphite  foil.  The  powder-filled  dies were  cold  pressed  at  approxi mately  50 MPa.  The  powders  were  sintered  using  field-assisted  sintering  (FAS: HPD 25-1, FCT 2000°C for  Sys teme, Rauenstein, Germany)  at  15 min  under  a  32 MPa  load. The controlled rates were 50°C/min. The load was applied to 1600°C and released on cooling  heating  and  cooling  during heating 1000°C.  to  Oxidation samples were cut with a wire electro-disinto 5 mm 9 5 mm 9 3 mm and 53 mm 9 3.5 mm 9 19-25 mm long  charge machine (EDM)  rectangles  (furnace heating)  5.0 mm rectangles with  a  centered  3 mm thick region of  reduced area (resistance heating).  The samples were polished using diamond slurry 1 lm finish.  to a  Polished  samples were  heated  by  a  zirconia-resis tance  furnace  (ZrF-25: Shinagawa Refractories, Tokyo,  Japan)  and direct  electrical  resistance. Macrographs of  the two tests and sample geometry are shown in Fig. 1.  The  furnace heating was  accomplished by a molybdeto 1100°C and a zirconia ele5°C/min. The  num disilicide pre-heater 2000°C at  ment  to  a  rate  of  samples  were  held  at  temperature  for  30 min.  Temperature  measurements  were  performed  using  a  single  color  pyrometer  focused  on  the  zirconia  element.  Samples  were  supported  on  a  zirconia  crucible. The  zirconia  crucibles  were  cut  from a  larger  crucible  (Advalue,  Tuscon, AZ; 10 mL Ca-stabilized ZrO2 ZrO2 and 4 ± 1% Ca). Direct electrical  crucible; 95%  resistance heat ing was  controlled  by  the  power  output  of  an AC  power  supply  across  the  sample  and  temperature was  read by a two-color pyrometer  (FMP2; FAR Associates,  Macedonia, OH)  that was  focused on the center of  the  reduced-thickness area. The samples were held in place  between two graphite  spacers by  tightening  set  screws  on  the  copper  electrodes. Ag  paint was  used  on  the  ends of  the samples  to improve electrical contact. Tem perature, current, and voltage data were recorded using  LabVIEW (National  Instruments, Austin, TX). Table I  lists  the oxidized samples with their heating conditions.  Oxidized samples were mounted in epoxy and pol ished in cross  section perpendicular  to the bottom (side  facing the crucible or notched side) of the sample to a 1 lm finish using diamond slurry. The microstructures  were  characterized  using  scanning  electron microcopy  (SEM, Quanta,  FEI, Hillsborough, OR)  along with  energy  dispersive  spectroscopy  (EDS,  Pegasus  4000;  EDAX, Mahwah, NJ)  for  elemental  analysis. All EDS  analysis was done using 15 kV accelerating voltage and  at  least a 100 s  live capture time.  Results  Single 2000°C Exposure  The heating profiles of  the HfB2-15 vol% SiC zir conia-resistance  furnace  heated  sample  (HS-F)  and  294  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  \\x0c', 'direct  electrical  resistance-heated  (HS-R)  sample  are  shown  in Fig. 2a  and  b, respectively. The maximum the HS-R sample was 2027°C  observed temperature of  using 82.5 V and 20.3 A (averaged over  the hold). The  oxidized  HS-F  sample  had  a  thicker  oxide  scale  (Fig. 3a) compared to the HS-R sample (Fig. 3b). The  HS-F sample was exposed to oxidizing temperatures  for  a greater length of time than the HS-R sample (6.5 h 1100°C compared ~4 min 800°C).  above  to  above  The oxide layers  labeled in Fig. 3a (HS-F) and Fig. 3b  (HS-R)  are  composed  of  (I)  a  SiO2-based  glass  that  penetrates  a HfO2-based skeleton;  (II)  a porous HfO2  scale;  and (III)  a SiC-depleted layer. The SiC-depleted  layer  is  defined  as HfB2 with a  reduced SiC content  (partially  oxidized SiC). The  average  total  oxide  scale  thickness measured from the top side of the HS-F sample is 660 ± 45 lm with the depleted layer comprising  53% of  the  scale. The  thickest  total oxide scale mea105 lm with 5% of  sured on the HS-R sample was  the total  scale consisting of  the depleted layer.  The oxide scales of HS-F samples possess a distinct  two-phase SiO2-based glass with the  less-pure  (less vis cous)  impurity-laden glass  rising  to the  surface of  the  oxide scale and the purer glass  found deeper within the  scale  (Fig. 4a).  A  two-phase  glass  found  in  furnace  heating has been described previously by  the  authors,  which was rities.15  shown to contain Al and Ca as major  impu In addition, HfSiO4  (with  a Ca  impurity)  is  found in the HS-F sample, but not  the HS-R sample.  The existence and absence of HfSiO4 was confirmed by  XRD. Figure 4b  is  an EDS  comparison of  the purer  (darker) and impure  (lighter) glasses  in the HS-F sam ple  along with the HfSiO4  and HfO2 phases.  In the  HS-R sample Al  impurities can be found randomly dis tributed  throughout  the  glassy  phase  (inset  Fig. 4c).  Figure 4d is  a  representative EDS spectra  of different  locations within the HS-R glass. There  is no hierarchy  to  the  concentration  of Al  in  the  glass  phase when  comparing  the  chemistry of  the glass  along the  length  of  the HS-R oxide scale.  Repeated 2000°C Exposure  An  advantage  of  the  direct  electrical  resistance  heating test  is  that  the sample can be exposed multiple  times  to the  same or different heating profiles. A sam2000°C twice  ple  (HS-Rr) was  heated  to  using  the  same heating profile as HS-R. The maximum observed 2030°C. The  temperature was  heating  profile  and  a  (a)  (c)  (b)  Fig. 1.  (a) Macrograph of direct electrical  resistance heating  apparatus. The temperature is read by means of a fiber optic  cable through a carbon tube (1)  that  leads  to the pyrometer.  Power is  supplied through the copper electrodes  (2)  to a heated  sample (3) gripped by carbon spacers  (4).  (b) Macrograph of  the  zirconia-resistance furnace showing the sample stand (A),  cylin drical zirconia heating element  (B), alumina insulation (C), and  molybdenum oxide insulation (D).  (c) Samples prepared for  direct electrical heating (top) and zirconia-resistance heating  (bottom) on the supporting zirconia crucible. The Ag paste  improves electrical  conduction.  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  295  \\x0c', '296  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  Table I.  List of  the Oxidized HfB2-15 vol% SiC Samples and their Heating Conditions  Sample ID  Test method  Max. observed temp. (°C)  Hold time (min)  Comments  HS-F  HS-R  HS-Rr  HS-R2  HS-Rf  Furnace  Self-heating  Self-heating  Self-heating  Self-heating  2000  2027  2030  2041  2325  30  1  1  2  0  —  —  Two 1 min holds  —  Heated to failure  Fig. 2.  Heating profiles of  the (a) HS-F (calculated) and (b)  HS-R (actual)  samples.  micrograph  of  the  resulting  oxide  scale  are  shown in  Figs. 5a and b. The oxide scale has a periodic structure  consisting  of  layers  of  SiO2  and HfO2 penetrated by  SiO2. For comparison, 2000°C  2 min.  for  to  observed  temperature  of  sured oxide scale of  a  sample  (HS-R2) was heated  (Fig. 5c) with 2040°C. The  a maximum  thickest mea found for  the HS-R sample  the HS-R2 sample was double that (217 vs 105 lm), and the  oxide  scale was not  composed of periodic  layers. The  lines.  the HS-F sample after oxidation at  Fig. 3. (a) Micrograph of 2000°C (b) Micrograph of the HS-R sample after oxidation at ~2000 °C. The oxide layers are (I) HfO2 penetrated by SiO2, (II) porous HfO2, and (III) depleted HfB2 layer. The approximate boundary between layers is shown by the dashed white  \\x0c', 'oxide  scale  formed near  the  center of  the  reduced area  on the HS-Rr  and HS-R2 samples were nonadherent.  Cracks were  also observed within the depleted layer of  the HS-F sample  and at  the  interface  between HfO2  and the SiC-depleted layer  (Fig. 3a).  In the HS-R sam ple,  fracture is observed between the depleted layer and  the HfO2  layer  at  the  center of  the  sample  (Fig. 3b),  whereas  adherent oxide  scales  exist near  the  end of  the  reduced area.  Temperatures Beyond 2000°C  The maximum temperature of  the direct  resistance  test  is  limited only by the available power and the sur vivability of  the  sample. A sample  (HS-Rf) was heated  to failure, where failure was defined as  the sample frac turing such that the electrical path was disrupted. The maximum observed temperature was 2325°C. A micro graph of  the cross  section of  the HS-Rf  sample (Fig. 6)  reveals extensive internal damage. Large pores are found  inside the sample whereas an oxide scale covers  the sur face. The bulk unoxidized material  from the  center of  the  sample  (inset Fig. 6) was  confirmed by EDS to be  SiC and HfB2. The microstructure  suggests  formation  of  a  liquid phase, which is 2347°C in  consistent with the calcu-SiC system.16  lated  eutectic  at  the HfB2  The oxide scale (inset Fig. 6) is composed of HfO2 penetrated by SiO2. Meng et al.13 similarly showed the -SiC sample failure of a ZrB2 at temperatures above 2300°C (2207°C eutectic temperature16), but did not  show any micrographs of  the interior microstructure.  Discussion  The  direct  comparison  of  the  zirconia-resistance -SiC  heated and direct electrical resistance-heated HfB2 2000°C provide  samples  at  insight  to  the  limitations  of  furnace  heating. Due  to  slower  heating  rates  and  contamination  from contact  between  the  sample  and  (a)  (b)  (c)  (d)  Fig. 4.  (a) Micrograph showing the different phases  found in the HS-F oxide scale (1) SiO2, (2) Si-Al-O, (3) HfSiO4 with Ca, and found in (a); (c) Micrograph showing the HfO2 (light) and Si-O-Al (dark) phases found the Si-Al-O phase corresponding to (c).  (4) HfO2; in the HS-R oxide scale;  (b) EDS corresponding to the phases  (d) representative EDS of  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  297  \\x0c', '298  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  (a)  (b)  (c)  Fig. 5.  (a) Heating profile of  sample HS-Rr.  (b) Micrograph of  Fig. 6.  Micrograph of the HS-Rf sample heated to 2325°C. The white-outlined inset shows HfB2 grains (labeled) and the eutectic SiC (dark)-HfB2 (light) structure found in the interior of the sample. The black-outlined inset is the oxide scale composed of HfO2 and SiO2 found on the exterior of the HS-Rf sample.  the oxide found in hottest region of the oxidized HS-Rr sample heated to 2000°C for 1 min two times. oxide found in hottest region of the oxidized HS-R2 sample heated to 2000°C for two minutes. The layers are the same as those found in the HS-R sample: (I) HfO2 penetrated by SiO2, (II) porous HfO2, and (III) depleted HfB2.  (c) Micrograph of  the  crucible,  the HS-F total oxide  scale  thickness  is greater  Fig. 7.  Micrograph showing a Si-Al-O impurity phase in the the HfB2-15 vol% SiC sample. The C signal in the EDS (inset) is from the carbon coating applied to the sample.  bulk of  than that observed in the HS-R sample. The difference  incongruently melting  silicate  and  thus  require  solid in  heating  rates  can  also  explain  the  observation  of  state  diffusion  to  form the  an  extra  kinetic  limitation on its  silicate phase adding formation.17 The  rapid  heating and cooling rates of  the HS-F sample presum ably do not allow for  the separation of glasses with dif HfSiO4  in  the HS-F  sample  but  not  sample. HfSiO4  is  only  stable  below  in the HS-R ~1726°C,17,18  therefore;  its  formation  in  the HS-F  sample  could  occur  during  slow cooling. HfO2  and  SiO2  form an  ferent viscosities or  for  the formation of HfSiO4.  \\x0c', 'In addition,  the  lack of Ca  impurity  in the  resis tance-heated sample suggests  the source of  the impurity  to  be  the CaO-stabilized  zirconia  crucible  or  zirconia  heating  element, whereas  the  presence  of Al  in  both  samples  implies  that  it  is  an inherent  impurity in the -SiC sample  starting powders. For  comparison,  a HfB2  was  heated  in  the  zirconia-resistance  furnace  using  a  Y2O3-stabilized  zirconia  crucible. The  glass  near  the  contact  region of  the sample and crucible was  found to  contain Si, Al, Ca, Y,  and O. Since  the  crucible was  reported to only contain 0.001% Ca,  the zirconia sam ple stand (Part A in Fig. 1) was  the likely source of Ca  in this  test. The HfB2  and SiC powders  are  reported  by the manufacturer  to contain 0.03% and 0.01% Al,  respectively. Figure 7 is a micrograph showing the SiC  grains with a pocket of  impurities  in the  as-processed  material. These areas can be found throughout  the sam ple  adjacent  to  SiC grains  and  are  shown  by  EDS  (inset)  to contain Si, Al, and O. The slow heating rates  and  contact  contamination  issues  of  the  zirconia  ele ment  furnace are not expected in hypersonic flight con ditions  and serve  to complicate  the  analysis of UHTC  oxidation resistance testing. -SiC sample is heated by direct elec When a HfB2  trical  resistance  through multiple heating  and  cooling  cycles,  spallation of  the oxide  scale is -HfO2  suggested by the  presence of  the  repeating SiO2  layers. Such lay ered oxide  structures have not  been reported for  fur nace-heated samples and was not observed in a sample  heated for  the same time (HS-R2) with a single heating  and cooling cycle. There  are  two sources of  stress dur ing oxidation that may lead to fracture during tempera ture  changes:  (i)  thermal  expansion mismatch and (ii)  volume  changes  associated with phase  transformations.  The  coefficient of  thermal  expansion (CTE)  of HfO2  depends on the impurity content and phase, but \\x006 K values are 5 9 10 to 7 9 10 for room temperature to 1250°C with purer HfO2 having lower values.19,20 Gasch et al.21 measured  typical  \\x006  \\x001  the CTE  of  pure  HfB2, pure SiC,  and a  combination of HfB2-20 vol%  SiC to find that \\x006 K the CTE of HfB2-20 vol% SiC was ~5 9 10 ~7 9 \\x006 K \\x001 at 1600°C and fell between the CTE values at room temperature and  \\x001  10  of pure HfB2  (higher)  and SiC (lower)  as  expected by  the rule of mixtures. The transformation of HfO2 from  monoclinic to tetragonal sion at 1642°C) or tetragonal to monoclinic cooling (10% conversion at 1710°C),18,22,23 3-3.5% volume  upon  heating  (10% conver during  is  accom panied  by  a  contraction/expansion  upon  heating/cooling.22,24  This  volume  expansion  could lead spallation of  the oxide scale.  As  the  absolute CTE and modulus of  the multi phase oxide  scale  are not known at  elevated tempera tures,  the main contributing  factor  to oxide  spallation  cannot  be  identified  definitively. However,  if  it  is  assumed that  the  volume  expansion upon phase  trans formation  is  isotropic  then  at minimum the  linear  expansion due  to phase  transformation would be 1%.  To achieve 2000°C to  greater  than  1% linear  expansion  from  room temperature when compared  to  the  bulk, the difference in CTE of the oxide scale and bulk \\x006 K \\x001. The would need to surpass ~4 9 10 reported -SiC and HfO2  range of CTE values for the bulk HfB2 ~2 9 10 \\x006 K  allow  for  a  \\x001  difference  between  the  CTE values, but  the difference could increase at higher  temperatures. Therefore,  it  is  possible  that  the  phase  transformation and CTE mismatch both contribute  to  spallation  of  the  oxide  scale. The  role  of CTE mis match and HfO2 phase  transformation on oxide  scale  adherence deserve further  study.  The limitation of the resistance heater was explored the sample was heated to failure above 2300°C. The  as  entire  sample was  soaked  at  the  elevated  temperature -SiC liquid  allowing for the formation phase inside the HS-Rf  of  the HfB2  sample. Furthermore,  the tem perature may  be  greater  in  the  interior  because  the  oxide scale will not be electrically conductive and is an  effective  thermal  insulator. Under  flight  conditions,  only  the outer  regions of  the  sample would be heated  and  the  high  thermal  conductivity  of  the  diboride  phase would  lead  to  a  temperature  gradient  through  the thickness of  the component. A temperature gradient  is  experienced along  the  length of  the direct  electrical  resistance  sample and can provide  insight  to oxide and  bulk microstructures over a temperature range.  Conclusion  Temperatures up to 2000°C can be  achieved in a  laboratory furnace; however,  these tests  suffer  from slow  heating profiles  and potential  interactions between fur nace materials and the sample being tested. The obser vation of Ca and HfSiO4  in the oxide  scale affects  the  glass  properties,  but  this  is  not  expected  in  a  flight  environment. The use of  resistance heating allows non contact  testing with a high heating profile. Features  like  fracture between the oxide  scale and the depleted layer  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  299  \\x0c', 'and Al  impurities  are  universal  observations  between  both heating tests  and require  further  investigation.  In  addition,  research to stabilize the tetragonal  transforma tion may aid in a more adherent  scale. Resistance heat ing may be  further utilized to  study multiple heating  profiles  and  test materials  for  scale  adherence.  The  resistance  testing  is  limited by  the uniform heating of  the  sample  that would not be  expected in a  real flight  environment. Further comparison of  test methods  such  as  laser heating, oxyacetylene  torch testing, or  scramjet  testing would  be  beneficial  to  understanding material  properties.  Acknowledgments  The authors  thank David Hart, AFRL Air Vehicles  Directorate,  for his assistance with the resistance heater  operation  and  Sindhura  Gangireddy,  University  of  Michigan,  for  the discussion involving microstructures -SiC to  of  resistively  heated  ZrB2  parallel  our  own  observations.  References  1.  J. R. Fenter, “Refractory Diborides 2 1-15 (1971).  as Engineering Materials,” SAMPE Q,  2.  J. W. Hinze, H. C. Tripp, and W. C. Graham, “High-Temperature Oxi dation Behavior of a HfB2 Plus 20 v/o SiC Composite,” Soc., 122 [9] 1249-1254 (1971).  J. Electrochem.  3. W. C. Tripp, H. H. Davis, and H. C. Graham, “Effect of an SiC Addition on the Oxidation of ZrB2,” Ceram. Bull., 52 [8] 612-616 (1973).  4. D. M. Van Wei, D. G. Drewry, D. E. King,  and C. M. Hudson,  “The  Hypersonic Environment: Required Operating Conditions Challenges,” J. Mater. Sci., 39 [19] 5915-5924 (2004).  and Design  5. T. H. Squire and J. Marschall, “Material Property Requirements  for Analy sis and Design of UHTC Components in Hypersonic J. Eur. Ceram. Soc, 30 [11] 2239-2251 (2010).  Applications,”  6. A. Bongiorno,  et al.  “A Perspective  on Modeling Materials  in Extreme  Environments: Oxidation of Ultrahigh-Temperature Ceramics,” MRS Bull., 31 [5] 410-418 (2006).  7. T. Sufang,  J. Deng, S. Wang, W. Liu, and K. Yang,  “Ablation Behaviors  of Ultra-High Temperature Ceramic Composites,” Mater. Sci. Eng. A, 465 [1-2] 1-7 (2007).  8. E. L. Corral  and L.  S. Walker,  “Improved Ablation Resistance  of C-C  Composites Using Zirconium Diboride Ceram. Soc, 30 [11] 2357-2364 (2010).  and  Boron Carbide,”  J.  Eur.  9.  J. Han, P. Hu, X. Zhang, S. Meng, ZrB2-SiC Composites  and W. Han,  “Oxidation-Resistant 799-806  at  2200  °C,” Comp.  Sci. Technol.,  68  (2008).  10. D. D.  Jayaseelan, H.  Jackson, E. Eakins, P. Brown, and W. E. Lee, “Laser  Modified Microstructures in ZrB2, ZrB2/SiC and ZrC,” Soc., 30 [11) 2279-2288 (2010).  J. Eur. Ceram.  11.  S. N. Karlsdottir and J. W. Halloran, “Rapid Oxidation Characterization of Ultra-High Temperature Ceramics,” J. Am. Ceram. Soc., 90 [10] 3233-  3238 (2007).  12. Z. Wang, Z. Wu, and G. Shi, “The Oxidation Behaviors of ZrB2-SiCZrC Ceramic,” Solid State Sci., 13 [3] 534-538 (2010).  13.  S. Meng, C. Liu, G. Liu, G. Bai, C. Xu,  and W. Xie,  “Mechanisms of  Material Failure for Fast Heating up at the Center of Ultra High Temperature Ceramic,” Solid State Sci., 12 [4] 527-531 (2010).  14. T. A. Parthasarathy, M. D. Petry, G.  Jefferson, M. K. Cinibulk, T. Ma thur, and M. R. Gruber, “Development of a Test  to Evaluate Aerothermal  Response of Materials  to Hypersonic Flow Using a Scramjet Wind Tun nel,” Int. J. Appl. Ceram. Technol., 8 [4] 832-847 (2011).  15. C. M. Carney, “Oxidation Resistance of Hafnium Diboride-Silicon Carbide From 1400 to 2000°C,” J. Mater. Sci., 44 [20] 5673-5681 (2009).  16.  L. Kaufman,  “Calculation of multicomponent  refractory  composite phase  diagrams,” NSWC TR 86-242 1 June 1986.  17.  J.-H. Lee, “Ternary Phase Analysis of Interfacial Silicates Grown in HfOx/ Si and Hf/SiO2/Si Systems,” Thin Solid Films, 472 [1-2] 317-322 (2005).  18.  S. V. Ushakov, A. Navrotsky, Y. Yang, S. Stemmer, K. Kukli, M. Ritali  et al.,  “Crystallization in Hafnia and Zirconia-Based Systems,” Phys. Sta tus. Solidi, 241 [10] 2268 (2004).  19.  S. L. Dole, O. Hunter, F. W. Calderwood,  and D.  J. Bray, “Microcrack[11-12] 486-490  ing  of Monolithic HfO2,”  J.  Am.  Ceram.  Soc,  61  (1978).  20.  S. R. Skaggs, “Zero and Low Coefficient of Thermal Expansion Polycrys talline Oxides,” LA-6918-MS September 1977.  21. M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S.  Johnson,  “Processing, Properties  and Arc  Jet Oxidation of Hafnium Diboride/Sili con Carbide Ultra High Temperature Ceramics,” 5925-5937 (2004).  J. Mater. Sci., 39 [19]  22. X. Luo, W. Zhou,  S. V. Ushakov, A. Navrotsky,  and A. A. Demkov,  “Monoclinic  to Tetragonal Transformations  in Hafnia  and Zirconia: A  Combined Calorimetric  and Density Functional Study,” Phys. Rev. B, 80  134119 (2009).  23. G. M. Wolten, “Diffusionless Phase Transformations nia,” J. Am. Ceram. Soc., 46 [9] 418-422 (1963).  in Zirconia and Haf 24.  L. Kaufman, E. V. Clougherty,  and J. B. Berkowitz-Mattuck,  “Oxidation  Characteristics of Hafnium and Zirconium Diboride,” Trans. Metall. Soc. AIME, 239 458-466 (1967).  300  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  \\x0c']"
},{
  "_id": 35,
  "PDF": "Densification and oxidation behavior of spark plasma sintered Hafnium Diboride-Hafnium Carbide composite.pdf",
  "Text": "['Ceramics International 46 (2020) 14625-14631  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Densification and oxidation behavior of spark plasma sintered Hafnium Diboride-Hafnium Carbide composite  T  Catalina Young, Cheng Zhang, Archana Loganathan, Pranjal Nautiyal, Benjamin Boesl, Arvind Agarwal∗  Plasma Forming Laboratory, Department of Mechanical and Materials Engineering, Florida International University, Miami, Florida, 33174, USA  A R T I C L E  I N F O  A B S T R A C T  Keywords: Hafnium diboride Hafnium carbide Spark plasma sintering Plasma jet oxidation Bimodal powder packing  This study reports on the sintering and oxidation behaviors of Hafnium Diboride (HfB2) and Hafnium Carbide (HfC) based ultra-high temperature ceramic (UHTC) composites. Pure Hafnium Diboride, Hafnium Carbide, and HfB2-50 vol% HfC composite were consolidated using spark plasma sintering (SPS) without the use of sintering aids. HfB2-HfC composite displayed a high sintered density of 98% as compared to 87% density of pure HfB2. The increased density of the HfB2-HfC composite is attributed to the bimodal powder distribution, which allowed for the smaller HfC particle to occupy the voids between the larger HfB2 particles resulting in improved packing efficiency. Moreover, the higher planar surface energy, for the (111) and the (200) planes for HfC, contributed to the smaller HfC particle being a significant driving force in the sintering process. Oxidation studies of each material were carried out by exposing them to a high-speed plasma jet in a temperature exceeding 2700 °C. The results of these studies show an improved oxidation resistance of HfB2-HfC composite by 54% and 70% over pure HfC and HfB2, respectively. The enhanced oxidation behavior is attributed to B2O3 filling in the porosity between the HfO2 scale and protecting the underlying material. This study provides a new alternative to improve the sinterability of UHTC diborides by introducing another UHTC (i.e., HfC) to form a fully dense composite without sintering aid and superior oxidation resistance.  1.  Introduction  Ultra-high temperature ceramics (UHTCs), including carbides, nitrides, and diborides of early transition metals, have been of increasing interest due to their extremely high melting point, high hardness, and elastic moduli. These materials are of particular importance for applications such as the nose and leading wing edge of hypersonic vehicle applications [1]. Most UHTCs cannot be used solely densified without using sintering aids, forming composites, or solid solutions [2-9]. The prime reasons for this situation is the low sinterability of the UHTCs due to the strong covalent bonds. Hafnium diboride (HfB2), among all the UHTCs, is recognized as one of the best oxidation-resistant ultrahigh temperature ceramics. 10,11,12 Upon oxidation, it forms porous HfO2 which serves as a scaffold filled with low melting point B2O3 [13]. Such combination results in dense oxidation layers that protect the underlying materials from further oxidation. However, B2O3 starts to evaporate around 1100 °C, which leaves the underlying materials vulnerable to oxidation [4]. Additionally, the sinterability of HfB2 is considerably poor; the sintered  samples are usually porous which provides an extra pathway for oxygen penetration and lead to further oxidation. Silicon carbide (SiC) is added as sintering aid as well as reinforcement for oxidation resistance in HfB2 [13]. Although the role of SiC in aiding sintering is yet to be fully understood, it is believed that silicon-based melts enable the liquid phase sintering which leads to higher densification in the final product. As for the oxidation resistance enhancement, the Si addition stabilizes the B2O3 by forming an amorphous borosilicate phase [10,11]. The latter phase is stable up to 1600 °C before it starts evaporating [14]. The drawbacks of the SiC addition are also glaring: the most studied HfB2SiC composition is HfB2-20 vol% SiC. Such a large amount of SiC will not only alter the properties of the HfB2 but also introduce unnecessary complex secondary phases to the composites. Although SiC can stabilize the B2O3 during oxidation up to 1600 °C, this temperature is still relatively low considering the desired temperature for UHTCs is usually higher than 2500 °C [10,14]. Thus, HfB2-SiC is still not the ideal UHTC composite. Hafnium carbide (HfC) is another ultrahigh temperature ceramic that has superior oxidation resistance. Unlike HfB2, which forms liquid  ∗ Corresponding author. E-mail address: agarwala@fiu.edu (A. Agarwal).  https://doi.org/10.1016/j.ceramint.2020.02.263 Received 31 December 2019; Received in revised form 18 February 2020; Accepted 27 February 2020  Available online 02 March 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'C. Young, et al.  Ceramics International 46 (2020) 14625-14631  phase oxide to seal the pores in the oxide scales, HfC forms a dense, crack-free oxycarbide layer when exposed to the oxygen that prevents further oxidation [3,5,13,15]. The oxycarbide layer is a metastable phase obtained from the oxygen absorption into HfC and substitute carbon atoms. HfC can absorb up to 30% of oxygen without transforming into HfO2. Zhang et al. sintered pure HfC without any sintering aid addition by spark plasma sintering (SPS) and achieved 98% dense pellet [5]. During plasma jet oxidation testing of pure HfC, it was observed that HfO2 formed due to the rapid speed of the oxygen entering the sample, and the oxycarbide layer formed was merely microns thick. Thus, leading to a thick oxide scale formation, of roughly 165 μm, on the surface of the sample [5-7]. In the present study, HfB250 vol% HfC (HfB2-HfC) UHTCs composites are sintered by spark plasma sintering and compared with pure HfB2 and HfC for their sintering and plasma jet oxidation behavior. It is well known that both HfB2 and HfC would form HfO2 during oxidation. It would be intuitive to mix these two ceramics such that the presence of B2O3 can preserve the oxycarbide layer from the HfC to further enhance the oxidation resistance. The addition of HfC to HfB2 would also improve the sinterability of the overall composite without introducing secondary phases. These two ceramics are expected to work in tandem to produce a superior UHTC composite with improved sinterability and oxidation resistance from the individual UHTC.  2. Experimental method  2.1. Material  Commercial powders were used to produce the hafnium ceramic composite system for the current study. Hafnium diboride powder was purchased from Sigma Aldrich, ST. LOUIS, MO, with a purity of 99 wt % size of 2.1 ± 0.83 μm. Hafnium carbide and the average particle powder was purchased from Materion, Tucson, AZ, with a purity of 99.5 wt % and the average particle size of 0.33 ± 0.08 μm.  2.2. Powder processing  For HfB2-50 vol% HfC, (here onward called HfB2-HfC) HfB2 and HfC were measured to be 50 vol % of mixed via ball milling in a high-energy vibratory ball mill machine (Across International LLC. Livingston, NJ, USA) for 1 h in a tungsten carbide milling jar and balls to obtain even particle distribution. The ball weight to powder weight ratio used was that of 1:3 to avoid degradation of particle size.  2.3. Spark plasma sintering  Compacts of pure HfC, pure HfB2, and HfB2-HfC, were fabricated via SPS (Model 10-4; GT Advanced Technologies, Santa Rosa, CA, USA). Powder (or mixture) was placed in 20 mm graphite die with two graphite punches, the die and punches were lined with graphite foil for easy removal. Each sample prepared was heated to a hold temperature of 1850 °C at 100 °C/min with 60 MPa of applied pressure under a vacuum of 4 Pa. Once the hold temperature was reached, the samples were held at 1850 °C for 10 min, then quickly cooled to room temperature. An optical pyrometer was used to measure the temperature of the samples during the SPS process via a hole cut into the die 5 mm away from the center wall of the die. Using the recorded ram displacement during the SPS run and the final sample height, initial sample height and instantaneous relative density during the process of SPS are calculated and obtained.  2.4. Characterization  Akishima, Tokyo, Japan). X-ray diffraction (XRD) of the powders and sintered pellets was conducted using a Siemens D-5000 X-ray diffractometer (Siemens AG, Berlin, Germany) at an operating voltage of 40 kV and 35 mA current, with the radiation source being CuKα. Density measurements were taken of each powder sample via a helium gas pycnometer (Accupyc 1340; Micrometrics Instrument Corporation, Norcross, GA, USA.). After sintering, graphite foil left behind from SPS was ground off the samples using diamond grit paper ranging from 125 to 6 μm and was metallographically polished using a 0.3 μm diamond suspension. XRD analysis was conducted on each sample after being polished to determine the phase retention and formation. Samples were then cut using a low-speed diamond saw and density measurements were taken via gas pycnometer. A comparison of theoretical and powder density versus measured density post-SPS was carried out to determine the densification of each compact. The cut samples were fractured by a hammer, and SEM and EDS analysis were conducted on the fracture surface to determine the microstructure and changes in composition after SPS. Grain size analysis was conducted using “Image J Fiji” software to measure grain growth from pure compacts to the mixed compact. Nanoindentation method was used to obtain elastic modulus, and nano-scale hardness of the sintered pellets via Hysitron Triboindenter TI900 (Hysitron INC, Minneapolis, MN, USA), with an 8000 μN applied load and Berkovich tip. The Oliver-Pharr method was used to compute elastic modulus. A minimum of fifteen indents was made for each sintered pellet.  2.5. Plasma jet oxidation testing  A Praxair SG-100 DC plasma gun was used for oxidation testing conditions to simulate reentry conditions of hypersonic vehicles where samples will experience both oxidation and ablation. Argon and Helium gases were used as the primary and secondary gases to form plasma, flowing at 56 slpm and 60 slpm, respectively. Gas flow velocity and temperature of the sample at the 75 mm distance from the t plasma gun were determined via Accuspray In-flight Particle Diagnostic Sensor (Tecnar Automation Ltd., QC, Canada). AlO-101 powder (Praxair Surface Technologies, Inc. Indianapolis, IN, USA) was used as spraying powder to measure, and the plasma jet was calculated to have a temperature exceeding 2700 °C and velocity of 330 m/s [7]. Each sample was held 75 mm away from the nozzle of the plasma gun to allow for oxygen exposure of the sample surface for 3 min. An in-house designed and fabricated fixture was used for the plasma jet oxidation tests [7]. This was done via a steel tube connected to a vacuum containing a coiled thermocouple to allow for the sample to be held without external mechanical forces applied via clamps, measurement of the backside temperature of the sample being tested, as well as allowing for direct exposure to the plasma jet [7].  2.6. Post-oxidation characterization  After plasma jet oxidation testing, XRD analysis was conducted on the top surface (facing plasma) of each sample to determine the formation of new phases. The top surface and cross-section morphologies of the oxide compacts were examined via SEM. The depth of the oxide layer was measured from the cross-sectional images using the “Image J Fiji” software.  3. Results & discussion  3.1. Effect of ball milling on powder  As-received UHTC powders and ball-milled HfB2-HfC were characterized via scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) using JEOL 6330F FE-SEM (JEOL Ltd.,  Fig. 1 shows SEM images of as-received HfC, as-received HfB2, and ball-milled HfB2-50 vol% HfC powder. HfB2 particle size is nearly seven times that of HfC, with a particle size of 2.1 μm and 0.33 μm, respectively. Tungsten carbide milling media was used to break down the  14626  \\x0c', 'C. Young, et al.  Ceramics International 46 (2020) 14625-14631  Fig. 2. Comparison of X-Ray diffraction patterns of ball-milled HfB2HfC, HfB2, and HfC powders, showing corresponding planes.  the sample by hindering crack propagation and increasing the flexural strength of the sample [16,17]. Fig. 2 shows the x-ray diffraction patters for Hf-based UHTC powders. Contamination by tungsten carbide milling media was not observed in the x-ray diffraction peaks.  3.2. Densification and sintering behavior during SPS  Table 1 summarizes the density and grain size of sintered HfC, HfB2, and HfB2-HfC pellets. The corresponding SEM images of the fracture surface are shown in Fig. 3. It can be observed that HfB2 has poor densification of 87.3% (Fig. 3b), whereas HfC and HfB2-HfC (Fig. 3c) pellets achieved a final density of 98% or more without the use of a sintering aid and no solid solution was observed (see Fig. 4). Fig. 3c shows a bimodal grain size distribution in dense HfB2-HfC pellet where large grains belong to HfB2 with finer grains of HfC. As can be seen in the in Table 1, there is a refinement in the microstructure of the grain size of HfB2 from the pure HfB2 compact to the HfB2-HfC compact by 34%, as well as a growth in HfC grain size from pure to mixed compact by 6%. The addition of HfC aids in the refinement of microstructure and leads to improved mechanical properties.  3.2.1. Sintering mechanism The densification curves for HfC, HfB2, and HfB2-HfC obtained from the ram displacement during SPS are shown in Fig. 5. Densification curves can be broken down into several empirical stages to aid in the understanding of the sintering and microstructure of HfB2-HfC. Stage 1 of SPS is the physical rearrangement of powder particles of all samples within the die as the load is applied. During stage 2, as the temperature is ramping up to the dwell temperature, the ductile-brittle transition temperature of HfB2, 1100 °C [4], is reached, and the plastic  Fig. 1. SEM images of a) as-received HfC powder, b) as-received HfB2 powder, and c) ball-milled HfB2 - 50 vol% HfC powder.  agglomerates and homogenously mix HfB2 and HfC powders as seen in Fig. 1c. It was noticed there was negligible change in particle size pre and post-mixing of the powders. This bimodal powder size distribution can be seen in Fig. 1c. Studies by Wonisch et al. and Khalil et al. have shown that bimodal powder size distributions have led to increased particle packing during sintering [16,17]. This increased particle packing has also been shown to improve the mechanical properties of  14627  \\x0c', 'C. Young, et al.  Table 1  Ceramics International 46 (2020) 14625-14631  A summary of powder size, sintered grain size, relative density, nano hardness, and elastic modulus of sintered HfC, HfB2, and HfB2-HfC pellets.  Compound  Relative Density  Powder Size (μm)  Grain Size (μm)  Elastics Modulus (GPa)  Nano Hardness (GPa)  HfC HfB2 HfB2-HfC  98.50% 87.30% 98.02%  0.33 ± 0.08 2.10 ± 0.83 0.33-2.10  0.73 ± 0.20 3.50 ± 1.17 0.78 ± 0.24  2.29 ± 0.70  300.4 ± 16.99 215.26 ± 63.95 354.5 ± 11.68  17.57 ± 0.88 14.35 ± 5.74 23.44 ± 2.13  Fig. 4. Comparison of X-Ray diffraction patterns of and HfC pellets, showing corresponding planes.  sintered HfB2-HfC, HfB2,  Fig. 3. SEM images showing fracture surfaces of sintered a) HfC, b) HfB2, and c) HfB2-HfC pellets.  Fig. 5. Empirical stages during SPS densification of HfC, HfB2, and HfB2-HfC.  14628  \\x0c', 'C. Young, et al.  Ceramics International 46 (2020) 14625-14631  Fig. 6. Selected portion of plotted current vs. HfB2HfC and HfB2.  temperature data from SPS of  deformation of HfB2 begins along its slip systems in both the pure HfB2 and HfB2-HfC samples, mentioned later. Stage 3 of SPS is the onset of consolidation of HfC, by self-diffusion and grain growth, which can be seen in both curves for HfC and HfB2-HfC. In stage 4, it can be seen the HfB2 compact begins consolidation into stage 5. During stage 5, it can be seen that the HfB2-HfC compact does not further consolidate, whereas the HfB2 compact continues to consolidate slowly with a significant grain growth. Comparing the densification curves in Fig. 5, it is seen that HfB2-HfC follows the consolidation trend of HfC until stage 4. This behavior suggests the sintering mechanism is dominated by HfC in HfB2-HfC composite [20]. The improved densification of HfB2-HfC is mainly attributed to bimodal powder distribution, which packs fine HfC between larger HfB2 particles. The high degree of interparticle packing (Figs. 1c and 2c) increases the particle contact between HfB2 and HfC. With more points of contact between powder particles, heating during SPS became more effective and uniform [19]. Fig. 6 is a plot of applied current vs. temperature during SPS for both pure HfB2 and HfB2-HfC samples. Since our SPS process is temperature-controlled, the amount of current applied reflects the effectiveness of the heating process. As shown in Fig. 6, the mixed sample always required a lower amount of current to achieve the same temperature as pure HfB2. This is due to the increased packing density, which helps in the effective hear transfer. On the contrary, in the case of pure HfB2 sintering, the heat cannot dissipate effectively due to pores from the lower packing density. As a result, more current was needed to compensate for the low efficiency in HfB2. Both planes with highest intensity in HfC (111) and HfB2 (0001) are close-packed planes filled with Hf atoms. Judging from the intimate interfaces between HfB2 and HfC after sintering (Fig. 3c), it is highly plausible that some localized diffusion occurred between these two planes and resulted in a higher densified sample. In addition to the localized diffusion, the hexagonal HfB2 particles undergo localized plastic deformation at the extreme processing conditions of high temperature and pressure during SPS, which further aids in the densification. The plastic deformation of HfB2 is due to the brittle to the ductile transition temperature of HfB2 at 1100 °C [4]. Evidence of this plastic deformation can be seen in the slip bands in the fracture surfaces of both the pure HfB2 and HfB2-HfC compacts in Fig. 7. These slip bands are due to the metal to metal bonding between the hafnium atoms on the {0001}, {1000}, and {1001} plane families. Due to the metallic character of Hf-Hf bonds, it is along these planes that slip occurs within the AlB2-type crystal structure, displayed here by HfB2 within the composite system [18].  Fig. 7. Comparison of slip lines HfB2HfC.  seen in the microstructure of a) HfB2 and b)  Fig.  8. Load-displacement noindentation.  curves  of  HfB2,  HfC,  and  HfB2-HfC  by  na 14629  \\x0c', 'C. Young, et al.  Ceramics International 46 (2020) 14625-14631  The effect of sintering on the nanohardness and elastic modulus of HfC, HfB2, and HfB2-HfC pellets can also be observed in Table 1 and Fig. 8. With the increased density and grain size refinement, HfB2-HfC composite showed the highest nanohardness and elastic modulus than those recorded of the constituent compounds. The low hardness and elastic modulus values of pure HfB2 are attributed to the high porosity. HfB2 also shows a wide scatter in the mechanical properties data which is due to the variation in the localized interaction zone of the nanoindenter probe with the porosity or sintered region. There was no sign of localized fracture from the load-displacement curve. A few curves for pure HfB2 showed a slight displacement at the peak load, which is attributed to localized slip, as observed in Fig. 7a.  3.3. Oxidation behaviors  Oxidation behaviors of sintered pure HfB2 and HfB2-HfC ceramic composite pellets were examined under high temperature and high gas flow speed to simulate real-world application conditions. The extreme conditions were simulated using a plasma torch, which can generate temperature up to 3000 °C and gas flow at sonic speed. Each pellet was exposed to the plasma torch for 180 s with a stand-off distance of 75 mm. The detailed experimental setup can refer to our previous work [7]. Oxidation behavior of a sintered pure HfC sample was taken from our last study, in which the HfC pellet was subjected to the same testing conditions as in the current work. After the oxidation tests, the oxide scales for each sample were measured from the cross-sectional images taken from SEM shown in Fig. 9. Unlike the HfC post oxidation sample which has a distinct oxide scale, the oxide scales for HfB2 and HfB2-HfC post oxidation samples are not distinctly discernible. To ensure the accuracy of the oxide thickness measurement, EDS analysis (results are not shown here) was performed on both samples and oxygen content was used to identify the oxide scale thickness. The thickness of oxide layers for HfC, HfB2, and HfB2HfC are 165, 250, and 75 μm, respectively. HfB2-HfC composites display the best oxidation resistance with an improvement of 54% over pure HfC and 70% over pure HfB2. The oxidation mechanisms for both pure HfC and HfB2 are well studied in previous literature [4,6,7,10-14,21,22]. In short, the resultant products from the HfC oxidation process are HfO2 and gases such as CO and CO2. Consequently, the oxide scale tends to be porous HfO2 as shown in Fig. 9a. In the case of HfB2, the oxidation products are HfO2 and B2O3. The melting point of B2O3 is ~400 °C [4,21-23] and it melts into a glassy phase during the oxidation process. The molten B2O3 fills the pores between HfO2 and results in a much denser oxide scale to protect the underlying material. Hence, the oxide scale in HfB2 would be thin and dense. In Fig. 9b, the sign of the molten glassy phase is presented as the dark and smooth region on the top surface of the oxide scale. The thickness of the HfB2 oxide scale is measured around 250 μm, which is significantly larger than that of the HfC oxide scale. This is contradictory to most of the literature [21,23] as HfB2 is expected to have better oxidation resistance as compared to HfC. To explain the contradiction, a top view SEM image of the HfB2 oxide scale is shown in Fig. 10a. The front side of the HfB2 oxide scale is relatively smooth. Numerous recrystallized grains indicate the presence of a significant amount of molten glassy phase formed during the oxidation process. However, pores (pointed by white arrows) are also visible. The size of these pores is similar to the pore size of sintered HfB2. This suggests that even though molten B2O3 was formed during the oxidation process, it is still not enough to seal the pores from the intrinsic HfB2 porosity. The interconnected pores in HfB2 provide a pathway for oxygen to penetrate a deeper layer of HfB2, which leads to more severe oxidation. Much improved oxidation resistance is observed in HfB2-HfC composite, as shown in Fig. 9c. The oxide scale of HfB2-HfC is dense and μm with only 75 signs of moderate melting (dark regions scatter through an oxide scale). A similar front side SEM image of the oxide scale is also provided in Fig. 10b. Compared to the front side surface of  Fig. 9. Comparison of oxide layers of each sample post 3-min plasma-jet oxidation testing.  the pure HfB2 oxide scale, the HfB2-HfC sample is denser and smoother. Less degree of melting is observed as compared to the pure HfB2 sample. However, the B2O3 was able to fill in more porosity between HfO2 and protect the underlying materials to near full densification. It is well accepted that Hf/ZrB220 vol% SiC is so far the best oxidation resistance UHTC composites [11,13,15,21,23]. SiC can not only improve the sinterability of the diborides but also stabilize the B2O3 and  14630  \\x0c', 'C. Young, et al.  Ceramics International 46 (2020) 14625-14631  influence the work reported in this paper.  Acknowledgment  Authors would like to thank Partnership for Research and Education Consortium in Ceramics and Polymers (PRE-CCAP) formed via the Department of Energy grant DE-NA0003865. The characterization facilities provided by Advanced Materials Engineering Research Institute (AMERI) at Florida International University are also acknowledged.  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2020.02.263.  References  [2]  [5]  [6]  [7]  [9]  [10]  [1] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. Zhou (Eds.), Ultra-high Temperature Ceramics: Materials for Extreme Environment Applications, John Wiley & Sons, 2014 Oct 10. F. Monteverde, Hot pressing of hafnium diboride aided by different sinter additives, J. Mater. Sci. 43 (3) (2008) 1002-1007. [3] D. Sciti, L. Silvestroni, A. Bellosi, Fabrication and properties of HfB2-MoSi2 composites produced by hot pressing and spark plasma sintering, J. Mater. Res. 21 (6) (2006) 1460-1466. [4] A.J. DeGregoria, Creep and Oxidation of Hafnium Diboride Based Ultra High Temperature Ceramics at 1500C, Air Force Institute of Technology WPAFB United States, 2015. C. Zhang, A. Gupta, S. Seal, B. Boesl, A. Agarwal, Solid solution synthesis of tantalum carbide‐hafnium carbide by spark plasma sintering, J. Am. Ceram. Soc. 100 (5) (2017) 1853-1862. C. Zhang, A. Loganathan, B. Boesl, A. Agarwal, Thermal analysis of tantalum carbide-hafnium carbide solid solutions from room temperature to 1400° C, Coatings 7 (8) (2017) 111. C. Zhang, B. Boesl, A. Agarwal, Oxidation resistance of tantalum carbide-hafnium carbide solid solutions under the extreme conditions of a plasma jet, Ceram. Int. 43 (17) (2017) 14798-14806. [8] R. Hassan, S. Omar, K. Balani, Solid Solutioning in ZrB2 with HfB2: effect on densification and oxidation resistance, Int. J. Refract. Metals Hard Mater. (84) (2019) 105041. B.C. Schulz, D. Butts, G.B. Thompson, Oxidation Behavior of vacuum plasmasprayed hafnium-tantalum nitrides, J. Mater. Res. 30 (19) (2015) 2949-2957. J.F. Justin, A. Jankowiak, Ultra high temperature ceramics: densification, properties, and thermal stability, AerospaceLab (3) (2011) 1. [11] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics, J. Mater. Sci. 39 (19) (2004) 5925-5937. [12] M. Gasch, S. Johnson, Physical characterization and arcjet oxidation of hafniumbased ultra-high temperature ceramics fabricated by hot pressing and field-assisted sintering, J. Eur. Ceram. Soc. 30 (11) (2010) 2337-2344. [13] R. Licheri, R. Orrù, C. Musa, A.M. Locci, G. Cao, Consolidation via spark plasma sintering of HfB2/SiC and HfB2/HfC/SiC composite powders obtained by self-propagating high-temperature synthesis, J. Alloys Compd. 478 (1-2) (2009) 572-578. J.F. Justin, Sintering and properties of Ultra High Temperature Ceramics for aerospace applications. ODAS 2013, (May 2013) PALAISEAU, France. [15] D.W. Ni, J.X. Liu, G.J. Zhang, Microstructure refinement and mechanical properties improvement of HfB2-SiC composites with the incorporation of HfC, J. Eur. Ceram. Soc. 32 (10) (2012) 2557-2563. [16] N.Z. Khalil, S.K. Vajpai, M. Ota, K. Ameyama, Effect of particle size distribution on SiC ceramic sinterability, Mater. Trans. 56 (11) (2015) 1827-1833. [17] A. Wonisch, T. Kraft, M. Moseler, H. Riedel, Effect of different particle size distributions on solid‐state sintering: a microscopic simulation approach, J. Am. Ceram. Soc. 92 (7) (2009) 1428-1434. [18] D. Ghosh, G. Subhash, N. Orlovskaya, Slip-line spacing in ZrB2-based ultrahightemperature ceramics, Scripta Mater. 62 (11) (2010) 839-842. S. Diouf, C. Menapace, A. Molinari, Study of effect of particle size on densification of copper during spark plasma sintering, Powder Metall. 55 (3) (2012) 228-234. [20] R.M. German, Prediction of sintered density for bimodal powder mixtures, Metal. Trans. A 23 (5) (1992) 1455-1465. C.B. Bargeron, R.C. Benson, R.W. Newman, A.N. Jette, T.E. Phillips, Oxidation mechanisms of hafnium carbide and hafnium diboride in the temperature range 1400 to 2100 C, Johns Hopkins APL Tech. Dig. 14 (1) (1993). T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, A model for the oxidation of ZrB2, HfB2 and TiB2, Acta Mater. 55 (17) (2007) 5999-6010. [23] R. Savino, M.D. Fumo, L. Silvestroni, D. Sciti, Arc-jet testing on HfB2 and HfC-based ultra-high temperature ceramic materials, J. Eur. Ceram. Soc. 28 (9) (2008) 1899-1907.  [19]  [21]  [22]  [14]  Fig. 10. Front side oxide scale of (a) HfB2 and (b) HfB2-HfC.  form a borosilicate glassy phase. However, the melting point of SiC itself is relatively low compared to those of UHTCs and would lead to premature failure in extreme conditions. Our study suggests a new route to improve the sinterability of the diborides by introducing another UHTC (i.e., HfC) with better sinterability. The second UHTC and bimodal particle size distribution can serve as a sintering aid to fully dense UHTC composites with a high melting point and superior oxidation resistance.  4. Conclusion  In the present work, a UHTC composite HfC-50 vol% HfB2 has been successfully sintered to near-full density. HfC and HfB2 are complementary to each other and result in a composite with improved sinterability, mechanical properties, and oxidation resistance. The improved sinterability is attributed to the bimodal powders added in the system, which leads to increased packing density and allows more effective and homogenous heating during the SPS process. Much improved densification leads to enhanced mechanical properties and plasma jet oxidation resistance over 54%. This work sheds new light on synthesizing advanced UHTCs composites without using sintering aids.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to  14631  \\x0c']"
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  "_id": 36,
  "PDF": "Densification of ZrB2-based composites and their mechanical and physical properties- A review.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 29 (2009) 995-1011  Review  Densification of ZrB2-based composites and their mechanical and physical properties: A review  Shu-Qi Guo  ∗  Composites and Coatings Center, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan  Received 21 August 2008; received in revised form 27 October 2008; accepted 4 November 2008  Available online 20 December 2008  Abstract  This study reviews densification behaviour, mechanical properties, thermal, and electrical conductivities of the ZrB2 ceramics and ZrB2 -based composites. Hot-pressing is the most commonly used densification method for the ZrB2 -based ceramics in historic studies. Recently, pressureless sintering, reactive hot pressing, and spark plasma sintering are being developed. Compositions with added carbides and disilicides displayed significant improvement of densification and made pressureless sintering possible at ≤2000 C. Reactive hot-pressing allows in situ synthesizing and densifying of ZrB2 -based composites. Spark plasma sintering displays a potential and attractive way to densify the ZrB2 ceramics and ZrB2 based composites without any additive. Young’s modulus can be described by a mixture rule and it decreased with porosity. Fracture toughness displayed in the ZrB2 -based composites is in the range of 2-6 MPa m1/2 . Fine-grained ZrB2 ceramics had strengths of a few hundred MPa, which increased with the additions of SiC and MoSi2 . The small second phase size and uniform distribution led to higher strengths. The addition of nano-sized SiC particles imparts a better oxidation resistance and improves the strength of post-oxidized ZrB2 -based ceramics. In addition, the ZrB2 -based composites showed high thermal and electrical conductivities, which decreased with temperature. These conductivities are sensitive to composition, microstructure and intergranular phase. The unique combinations of mechanical and physical properties make the ZrB2 -based composites attractive candidates for high-temperature thermomechanical structural applications. © 2008 Elsevier Ltd. All rights reserved.     Keywords: Zirconium diborides; Densification; Mechanical properties; Thermal and electrical conductivities  Contents  1. 2.  3.  4.  5.  Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Densification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Hot-pressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Spark plasma sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. Reactive hot-pressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4. Pressureless sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mechanical behaviours . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Young’s modulus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Fracture toughness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Flexural strength . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Physical behaviours . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Thermal conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Electrical conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .  996 996 996 998 999 1000 1001 1001 1003 1004 1006 1007 1008 1009 1009  ∗  Tel.: +81 29 859 2223; fax: +81 29 859 2401.  E-mail address: GUO.Shuqi@nims.go.jp.  0955-2219/$ - see front matter © 2008 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2008.11.008  \\x0c', '996  1.  Introduction  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  is necessary to understand the relation of performance to processing, compositions, and microstructure in order to produce an ultra-high temperature ZrB2 -based composite with superior performance. In this article, the densification behaviour, Young’s modulus, fracture toughness, flexural strength, thermal and electrical conductivities of the ZrB2 ceramics and ZrB2 -based composites are reviewed. The emphasis is directed toward presenting recent advances and providing an evaluation of studies of the ZrB2 -based ceramics materials.  2. Densification  Recently, ZrB2 ceramics and ZrB2 -based composites have been densified by the various methods, including hot pressing (HP), spark plasma sintering (SPS), reactive hot pressing (RHP), and pressureless sintering (PS). Historically, HP has been the dominant method in densification studies. Later, SPS, RHP, and PS processes evolved as the most common densification methods. This section focuses on these four densification processes, as well as on the effects of additives, such as carbides, nitrides and disilicides, on the densification behaviour of zirconium diborides.  2.1. Hot-pressing (HP)           The densification of ZrB2 powder generally requires very high temperatures,22 owing to the covalent character of the bonding as well as to its low volume and grain boundary diffusion rates. Typically, HP of ZrB2 required a temperature of 2100 C or above and moderate pressure (20-30 MPa),23-25 or temperatures (1800 lower C) and extremely high pressures (>800 MPa).26,27 These studies23-27 showed that densification of ZrB2 is a diffusion-controlled rate process. Results of HP studies on commercially available ZrB2 powders are summarized in Table 1, which includes details of the starting powders, sinterIt has been found that HP of coarse ZrB2 powder (d ≈ 20 \\u242em) ing additives, HP conditions, mixing method and final densities. at 2000 C with a pressure of 20 MPa achieved only a relative density of 73%,24 whereas the relative density of 91% was obtained for a finer ZrB2 powder (d ≈ 2.1 \\u242em) under the same HP conditions.25 Furthermore, the attrition-milled ZrB2 powder, with average particle size of d ≤ 0.5 \\u242em, required HP at 1900 C and 32 MPa for 45 min to achieve full density.11 The ticle size from microns (d ≈ 2.1 \\u242em) to submicrons (d ≤ 0.5 \\u242em) lower HP temperature was attributed to reduction of starting parby attrition-milling. Oxygen impurities (B2O3 and ZrO2 ) present on the starting powder surfaces have been shown to inhibit densification and to promote grain growth in the non-oxide ceramic systems. A study28 in TiB2 suggested that the total oxygen content must be less than 0.5 wt% to achieve full density. Recently, metal lic Ni,29,30 SiC,10-16,25 Si3N4 ,31,32 AlN,33,34 HfN or ZrN,35,36     have been added to ZrB2 , producing an intergranular secondary phase and/or reducing oxygen content, both of which assists in the densification of ZrB2 . Silicon carbide is the most common additive for ZrB2 or HfB2 ceramics. The addition of SiC                          Structural materials for use in high-temperature oxidizing environments are presently limited to SiC, Si3N4 , oxide ceramics, and composites of these materials. Silicon-based ceramics are oxidation resistant up to 1600 C, due to the formation of a protective SiO2 surface film.1,2 Although SiO2 is an excellent oxidation barrier at temperatures below 1600 C, and above this temperature it begins to soften dramatically and develops substantial vapour pressure.1,2 Therefore, the use temperature of the silicon-based ceramics is limited to 1600 C by their thermal stability in an oxidizing atmosphere. In addition, there are a relatively few refractory oxides that are stable in an oxidizing environment at or above 2000 C. Among these oxides, zirconia (ZrO2 ) and hafnia (HfO2 ) typically have the highest melting points, 2700 C and 2800 C,1,3 respectively. Although they are inert chemically, they appear to remain susceptible to thermal shock, and exhibit high creep rates and phase transition at higher temperatures.1,3,4 Therefore, the development of structural materials for use in oxidizing environments at temperatures above 1600 C is of great engineering importance. Ceramics based on the transition metal borides, nitrides, and carbides have extremely high melting points (>2500 C) and are referred to as ultra-high temperature ceramics.1,5 Within the family of transition metal ultra-high temperature ceramics, diborides such as ZrB2 and HfB2 have unique combinations of mechanical and physical properties, including high melting points (>3000 C), high thermal and electrical conductivities, chemical inertness against molten metals, and great thermal shock resistance.1,5,6 Thus, although carbides typically have the highest melting points (>3500 C), the diborides ZrB2 and HfB2 are more attractive candidates for high-temperature thermomechanical structural applications at a temperature ≥2000 Potential applications for the diborides include thermal protective structures for leading-edge parts on hypersonic re-entry space vehicles,1,7 propulsion systems,1,7 furnace elements,8 refractory crucibles,8 and plasma-arc electrodes.8,9 In particular, ZrB2 has the lowest theoretical density among the ultra-high temperature ceramics, which makes it an attractive material for aerospace applications.1,5,7 However, the use of the single-phase material for high-temperature structural applications is limited by its poor oxidation and ablation resistance, as well as its poor damage tolerance. The composite approach has been successfully adopted in order to improve the densification, mechanical properties, physical properties, as well as the oxidation and ablation resistance of the ZrB2 ceramics.10-21 For example, the addition of 20 vol% fine SiC (d ≈ 0.5 \\u242em) increased the strength of ZrB2 to over 1 GPa.11 The ZrB2 -MoSi2 composites consolidated by spark plasma sintering (SPS) can retain their room temperature strength (650 MPa) to at least 1200 C.18 The addition of ZrSi2 reduced the densification temperature of ZrB2 to below 1550 C, as well as increasing its thermal and electrical conductivities.20,21 Obviously, the mechanical and physical properties of the ZrB2 -based composites are closely linked with the densification process, compositions, starting powder, microstructure, and intergranular second phase. Therefore, it  C.1,5                 \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  997  Table 1  Starting powder size, hot pressing conditions and final densities of the hot-pressed ZrB2 ceramics and ZrB2 -based composites with various additives.  Compositions (vol%)  Particles size (\\u242em)  Remarks  Hot-pressing conditions  Final density (%)  References  ZrB2  SiC or MoSi2 or ZrSi2  ZrB2 ZrB2 ZrB2 ZrB2 ZrB2 ZrB2 -2.5 wt% Si3N4 ZrB2 -5Si3N4 ZrB2 -4.6AlN ZrB2 -15SiC-4.5ZrN ZrB2 -37.5HfB2 -19.5SiC-3HfN ZrB2 -5.7SiC ZrB2 -22.4SiC ZrB2 -22.4SiC ZrB2 -30SiC ZrB2 -30SiC ZrB2 -16(SiC + C)  20  Ball-milled  2000                                   C/20 min/20 MPa  73  24  2.1  Ball-milled  2000  C/60 min/30 MPa  91  25  2  Attrition-milled  1900  C/45 min/32 MPa  99.8  11  2  Ball-milled  1650  C/20 min/60 MPa  71.6  13  5-10  Ball-milled  1800  C/60 min/20 MPa  78  37  0.1-8  Ball-milled  1700  C/15 min/30 MPa  98  31  2  Ball-milled  1700  C/15 min/30 MPa  98  32  0.1-8  Ball-milled  1850  C/30 min/30 MPa  92  33  d90 = 4-6 2  Ball-milled  1900  C/5 min/50 MPa  99  36  Ball-milled  1900  C/30 min/50 MPa  >99.9  35  2  1.7  Ball-milled  1650  C/120 min/60 MPa  81.6  13  2  1.7  Ball-milled  1650  C/120 min/60 MPa  97.9  13  2  0.04  Ball-milled  1650  C/120 min/60 MPa  99.6  13  6  10  Attrition-milled  1900  C/45 min/32 MPa  97.4  14  6  0.7  Attrition-milled  1900  C/45 min/32 MPa  98.7  14  5-10  Polycarbosilane  (PCS)  Ball-milled  1800  C/60 min/20 MPa  100  37  ZrB2 -20MoSi2 ZrB2 -20MoSi2 ZrB2 -20ZrSi2  2  2.8  Ball-milled  1800         C/5 min/30 MPa  98.1  17  2.1  3.1  Ball-milled  1800  C/30 min/30 MPa  99.8  19  2.1  2.5  Ball-milled  1400  C/30 min/30 MPa  +1550     C/15 min/30 MPa  99.1  21  ZrB2 -20MoSi2  20  3-5  Ball-milled  2000     C/20 min/20 MPa  95  24  improved the sinterability, inhibited grain growth and increased the oxidation and ablation resistance of ZrB2 and HfB2 ceramics as well.10-16,25 Monteverde12 showed that ZrB2 with 10 vol% ultra-fine SiC (d90 = 0.8 \\u242em) achieved full density by HP at 1900 C and 40 MPa for 20 min in vacuo. The attrition-milled (d ≈ 0.5 \\u242em) ZrB2 -30 vol% SiC mixture powders could be hot-pressed at 1900 C to a relative density exceeding 97%.14 average particles sizes ranging from 40 nm to 0.6 \\u242em, sharply Furthermore, the addition of 22.4 vol% nano-sized SiC, with reduced the HP temperature necessary to achieve full density to C (pressed for 120 min at 60 MPa).13 The improvement of 1650 densification upon addition of SiC was attributed to the formation of intergranular liquid phases during hot-pressing, assisting in densification at lower temperatures.12,13 An early study29 in 4 wt% Ni-containing ZrB2 showed that the presence of the liquid phase not only favours ZrB2 particle rearrangement but also enhances mass transfer kinetics. However, the improvement of densification upon addition of ultra-fine SiC is effective only for a uniformly dispersed SiC-ZrB2 system.25 The agglomeration of the ultra-fine SiC particles led to reduced improvement in densification of ZrB2 , even with nano-sized SiC particles.25 Recently, a polycarbosilane (PCS) was used as a source of SiC and C because the pyrolyzed PCS can crystallize and convert to ␤-SiC and amorphous C at 1000 C or above.37 HP of the PCScoated ZrB2 powder required a reduced temperature of 1800 C (pressed for 60 min at 20 MPa) to achieve a full density when SiC content was ≥16 vol%. For comparison, the ZrB2 powder without PCS coating was highly porous with a relative density of 78% under the same HP conditions. Nitrides are other effective additives for improving sinterability or enhancing densification of ZrB2 . The main reason for                 incorporating nitrides as additives is the propensity of nitrides to consume the oxygen-bearing species on the diboride powder surfaces. The reduction of oxygen results in higher boron activity, which is one of the conditions favouring lattice diffusion and, therefore, densification.28 The addition of ≥2.5 wt% Si3N4 results in almost fully dense ZrB2 (RD: 98%) after compactions at 1700 C and 30 MPa for 15 min.31,32 Some grain boundary phases, including BN, ZrO2 , ZrSi2 , and borosilicate glassy phase, were confirmed to be present in pockets at multiple-grains junctions for the ZrB2 -based ceramics with Si3N4 , the result of a reaction of an oxide impurity with Si3N4 . That reaction results in elimination or decrease in the oxide impurity on the ZrB2 particles surfaces, thereby promoting densification.31,32 Similar to Si3N4 , the primary effect of an AlN additive is the depletion of the ZrB2 particles from the outer oxide layer that prevents the formation of highly dense compacts.33 Compared to AlN and Si3N4 additives, the ZrN and HfN showed the unexpected advantage of limiting undesirable secondary phases that eventually become detrimental to high temperature stability. It has been reported that HP of 3 vol% HfN-HfB2-SiC35 or 4.5 vol% ZrN-ZrB2 -SiC36 required 1900 C for 30 min at a pressure of 40 MPa. The resulting composites showed a fine and homogenous microstructure with secondary phases such as M(C, N), MO2 (M = Zr and/or Hf) and BN. The formation of the secondary phases during sintering was traced back to the interactions among ZrN or HfN, carbon, and oxides such as B2O3 and ZrO2 or HfO2 . These interactions accelerated the densification of ZrB2 or HfB2 ceramics by reducing the oxygen content on the starting powder surfaces. The resulting intergranular secondary phases possess higher refractoriness than those made with AlN or Si3N4 additives.        \\x0c', '998  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011        Transition metal disilicides have been found to be an alternative and effective sintering additive because they improve sinterability and increase oxidation resistance of ZrB2 ceramics as well. In the early 1970s, Kinoshita et al.24 systematically investigated densification behaviour of ZrB2 -based composites with MoSi2 . They found that MoSi2 significantly improved sinwas obtained for ≥20 vol% MoSi2 -containing ZrB2 powder at terability of ZrB2 powder and a relative density exceeding 95% 2000 C and 20 MPa for 20 min. Recently, it has been reported that HP of the ZrB2 -based ceramics with MoSi2 required only a temperature ≤1800 C.17-19 The higher sintering temperature required in the earlier study resulted from larger ZrB2 (d ≈ 20 \\u242em) and MoSi2 (d ≈ 5 \\u242em) particles. More recently, Guo et al.21 found that addition of 10-40 vol% ZrSi2 could further lower the densification temperature of ZrB2 to 1550 C or below. Furthermore, fully dense ZrB2-ZrSi2 composites with a fine and homogeneous microstructure, using a two-step HP process, which consisted of a first stage at 1400 C for 30 min and a second stage at 1550 C for 15 min at a pressure of 30 MPa were obtained. Thus, the disilicides of the transition metals are potential additives for lowering the sintering temperature of the ZrB2 -based ceramics. Consequently, it is possible to further lower the sintering temperature of ZrB2 -based ceramics by selecting appropriate disilicides of transition metals. Improvement of densification, resulting from addition of disilicides, is attributed to two major causes. One is formation of an intergranular Si-O-B liquid phase between MoSi2 or ZrSi2 and ZrB2 particles due to the interaction of SiO2 and B2O3 that occurs on the surfaces of particles.38 Another is the ductile deformation of MoSi2 or ZrSi2 particles at high temperature (>800 This deformation could force soft MoSi2 or ZrSi2 particles to fill in the voids left by the ZrB2 skeleton under pressure during sintering, thus improving densification.19,21  C).39,40              2.2.  Spark plasma sintering (SPS)  SPS is one of the most recent advanced processing techniques developed for densifying ceramic materials.41,42 Although SPS is similar to HP, in place of indirect heating, the applied electrical field heats the die and the powder compact. One advantage of using SPS is to enhance densification of poorly sinterable ceramics, by simultaneously applying a uniaxial load and a direct or pulsed electric current to a powder compact. Another advantage is that the grain growth of starting materials is restricted, since a considerable shorter sintering time (within minutes) is needed compared to HP or hot isostatic pressing (HIP), thereby retaining the fine and homogeneous grains. Previous investigations43-45 of compaction of oxide, nitride, and carbide ceramics produced by SPS have shown that the sintering time, heating rate and sintering temperature are the important factors controlling fine-grained microstructure and densification. In particular, the selection of the sintering temperature is critical for the development of the optical microstructure. Recent studies have shown that SPS enhanced densification and refined microstructure of ZrB2 -based ceramics can be achieved in very short processing cycles.18,46,47 Table 2 summarizes SPS conditions, final density, and grain size of                    ZrB2 -based ceramics produced by SPS. Medri et al.46 showed that 60ZrB2-30ZrC-10SiC (vol%) composition could be sintered to a relative density of 96% at 2100 C and 30 MPa for 2 min. Grain size measurement indicated that the grain growth (maximum grain size: 3 \\u242em) was inhibited during SPS. Recently, the various ZrB2 -ZrC-SiC compositions could be sintered to the fully dense compacts with fine and homogenous microstructure at 1950 C and 30 MPa for 2 min, by using the SPS technique.48 This discrepancy in the sintering temperature is probably associated with starting powder size and SPS conditions. In addition, extending soaking time from 3 min to 5 min can produce fully dense ZrB2-ZrC-SiC composites at a lower temperature (1900 C).34 Furthermore, addition of 5 wt% AlN results in complete densification at 1850 C and 30 MPa for 5 min, but addition of 5 wt% Si3N4 still required a temperature of 1900 C. The discrepancy in densification temperature due to additions of AlN and Si3N4 is likely attributable to a lower onset temperature of densification and a faster shrinking rate for the AlN addition as compared to the Si3N4 addition.34 For the ZrB2 -15 vol% MoSi2 ,18 however, the density and grain size measurements of the compacts consolidated at 1750 C showed that SPS was not superior to HP. Soaking time and total sintering time were noticeably shorter for SPS (7-24 min) than for HP (20-140 min). Guo et al.49 systematically investigated the densification behaviour and grain growth of ZrB2 ceramics produced by SPS. They found that densification and grain size of the sintered ZrB2 compacts were strongly dependent on the selection of sintering temperature, holding time, as well as the heating rate. It was possible to obtain the almost fully dense ZrB2 ceramics with a fine and homogeneous microstructure by selecting the appropriate sintering parameters (Fig. 1(a)). Without sintering at temperatures ≥2100 additives, full density has, historically, been achieved only by HP C.23,24 Densification and grain growth occurred simultaneously during the sintering. As a result, it was difficult to obtain a full density ZrB2 compact. SPS of ZrB2 ceramics required a temperature of 1900 C, a holding time of 3 min, and a heating rate of 200 C/min or above. Increasing the sintering temperature to 1950 C or extending the holding time to 10 min or above, as well as lowering the heating rate below 200 C/min, led to coarsening of the grain size (typical example, Fig. 1(b)). In addition, SPS has been used for applications in other transition metal diborides, such as HfB2 -SiC,18 TiB2 -WB2-CrB2 ,50 TiC-TiB2 ,51 HfB2-MoSi2 ,52 and HfC and HfB2 -based composites with MoSi2 additives.53 The enhanced densification resulting from SPS was attributable to mass transfer processes, which are significantly enhanced in the process, effectively promoting densification. The mechanism in SPS technique that enhanced densification - mainly whether or how an electric discharge is involved in accelerating the densification and grain growth - is still the subject of intense debate. However, we suggest that enhancement is most probably due to (i) an efficient heat transfer; (ii) the use of comparatively high pressure; (iii) the presence of an electrical field (use of DC pulses); and (iv) the presence of local spark discharges generated between the powders under high-energy electrical pulses.                \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  999  Table 2  Spark plasma sintering conditions, final density, and grain size of ZrB2 ceramics and ZrB2 or HfB2 -based composites produced by an SPS process.  Compositions  Heating rate  SPS conditions  Final density (%)  Grain size (\\u242em)  References  ZrB2 -30ZrC-10SiC (vol%) HfB2 -30SiC (vol%) ZrB2 -15MoSi2 (vol%) (15-70)ZrB2 -(15-70)ZrC-(15-50)SiC(mol.%) 26.18ZrB2 -59.87ZrC-13.95SiC(wt%) 24.17ZrB2 -56.88ZrC-13.25SiC-5AlN(wt%) 24.17ZrB2 -56.88ZrC-13.25SiC-5AlN(wt%) ZrB2 ZrB2 ZrB2 HfB2 -(1-9)MoSi2 (vol%) HfC-9MoSi2 (vol%)  100         C/min  2100                           C/2 min/30 MPa/vacuum  96  3  46  100  C/min  2100  C/2 min/30 MPa/vacuum  100  2 1.4 -  46  400 100 C/min 400 400 400 300 300 300 100 C/min  1750  C/7 min/30 MPa/vacuum  97.7  18                 C/min  1950  C/2 min/50 MPa/Ar  >98  48  C/min  1900  C/5 min/30 MPa/Ar  99.9  -  34  C/min  1850  C/5 min/30 MPa/Ar  99.5  -  34  C/min  1900  C/5 min/30 MPa/Ar  100  -  34  C/min  1900  C/3 min/50 MPa/vacuum  97.6 80 98  5.1  49  C/min  1900  C/10 min/50 MPa/vacuum  10  49  C/min  1950  C/3 min/50 MPa/vacuum  19  49       1750  C/3 min/100 MPa/vacuum  >97  1  53  100  C/min  1750  C/3 min/100 MPa/vacuum  99  0.8  53  2.3. Reactive hot-pressing (RHP)  The use of metallic and ceramic additives during HP could reduce temperature of densification and also inhibit grain growth in ZrB2 ceramics. However, the significant decreases in the strength at temperatures above 1200 C that result from softening of intergranular amorphous phase at elevated temperatures has been reported for ZrB2 -based composites with SiC.18,27,29,30,32 RHP has been identified as a potential route     Fig. 1. FE-SEM backscattered electron image of the ZrB2 ceramics consolidated heating rate of 200 by SPS at 1900 C for different holding time of (a) 3 min and (b) 10 min with C/min under a pressure of 50 MPa in vacuum.49        to produce ZrB2 ceramics with low impurity levels and high density at a lower temperature. There are two processes that occur in RHP, in situ reaction of precursor powders and densification, which must be completed simultaneously during heating and subsequent holding. Recently, RHP has been used to produce ZrB2 and/or HfB2 dense compacts by using Zr and/or Hf and B precursors as well as to fabricate the ZrB2 -based composites with SiC and/or ZrC by using Zr, Si and B4C precursors. Table 3 summarizes RHP sintering conditions, precursors, grain size and final density of the ZrB2 ceramics and ZrB2 -based ceramics fabricated by RHP. Chamberlain et al.54 have employed slow heating (1 C/min) and extended isothermal holds at an extremely low temperature (6 h at 600 C) to react fine powders of Zr and B without ignition of self-propagating high-temperature synthesis (SHS) reaction. When the samples were heated to 1650 C in an argon atmosphere, an applied external pressure of 40 MPa produced an almost fully dense, nano-sized ZrB2 compact. Raising the temperature to 1700 C increased the density to 99%, however, the ZrB2 grains were significantly coarsened. The grain size meaC was 1.5 \\u242em, sured in the sample densified at 1800 larger by a factor of 3 than that at 1650 C. In contrast, HP of commercially available micron-sized ZrB2 powders (d ≈ 2.1 \\u242em) at 2000 C and 20 MPa for 60 min achieved only a relative density of 91%,25 with an average grain size of 6.1 \\u242em. The improvement of densification by RHP was attributed to the formation of nano-sized ZrB2 particles during the reactive process because the fine crystalline size should enhance the driving force for densification when the densification is driven by minimization of the surface free energy. Another application of RHP is to produce ZrB2 -based composites with SiC and/or ZrC, using Zr, Si, and B4C powders as precursors. Zhang et al.55 used RHP to fabricate ZrB2 -beased composites with SiC, by reacting Zr, Si and B4C at 1800 C where the following reaction is thermodynamically favourable: 2Zr + Si + B4C = 2ZrB2 + SiC (1) The relative density of 98% was obtained by RHP of Zr, B and B4C powder mixtures at 1900 C and 30 MPa for 60 min. Later, Wu et al.56 also successfully consolidated (RD: 97%) by RHP of Zr, Si ZrB2-SiC-ZrC composites                             \\x0c', '1000  Table 3  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  Reactive hot-pressing conditions, precursors, grain size and final density of ZrB2 or HfB2 ceramics and ZrB2 or HfB2 -based composites fabricated by a RHP method.  Materials  Precursors  Remarks  HP or SPS Processing conditions  Final density (%)  Grain size (\\u242em)  References  ZrB2 HfB2 HfB2 ZrB2 -SiC ZrB2 -SiC-ZrC ZrB2 -SiC ZrB2 -SiC HfB2 -SiC  Zr, B  Hf, B  HfB2 Zr, Si, B4C Zr, Si, B4C Zr, Si, B4C Zr, Si, B4C Hf, Si, B4C  Attrition-milled  None  None  Ball-milled  None  Ball-milled  Ball-milled  Ball-milled                   1650  1700  1900  1900  1800  1450  1890  1900  C/30 min/40 MPa/Ar (HP)  C/10 min/95 MPa/vacuum (SPS)  C/10 min/95 MPa/vacuum (SPS)  C/60 min/30 MPa/Ar (HP)  C/60 min/20 MPa/Ar (HP)  C/3 min/30 MPa/vacuum (SPS)  C/10 min/30 MPa/vacuum (HP)  C/60 min/50 MPa/vacuum (HP)  >95 98 87 96.7  96.8 98.5 100  100  ZrB2  SiC  0.5  -  -  3-10  -  <5  2  3  -  -  -  <3  -  <1  1  1  54  59  59  55  56  60  57  58           and B4C powders at 1800 C and 20 MPa for 60 min in an argon atmosphere. They found that the reactions for producing ZrB2 , ZrC and SiC were not simultaneously induced during the sintering, but in steps. ZrB2 and ZrC were first formed by the reaction of B4C with Zr at a low temperature, then SiC was produced by reaction of Si with ZrC and the residual B4C at a higher temperature. In addition, Zimmermann et al.57 found that excess B4C and Si were necessary in the ZrH2 -B4C-Si system for obtaining ZrB2-SiC composites without oxide impurity as well as for avoiding grain coarsening during the sintering process. They showed that RHP produced ZrB2 -based and Si, with an average ZrB2 grain size of 2 \\u242em and a SiC ceramics with 27 vol% SiC in the presence of excess B4C particle size of 1 \\u242em. For comparison, in a stoichiometric ZrH2 , B4C and Si mixture, the reaction resulted in 25 vol% SiC-containing ZrB2 , accompanied by traces of ZrC and ZrO2 , mixture. The grain size appeared to increase from 2 \\u242em to as a result of deficiency of available boron in the reaction 5 \\u242em for the ZrB2 and from 1 \\u242em to 3.5 \\u242em for the SiC. Similarly, RHP is also used to produce the HfB2 -based Monteverde58 composites. obtained a fully densified HfB2-22 vol% SiC-6 vol% HfC at 1900 C and 50 MPa for 60 min through reaction of a mixture of Hf, B4C and Si powders. Recently, RHP has also been utilized to produce HfB2 by reaction of Hf and B powders at a low temperature, by using SPS. An almost fully densified HfB2 compact was achieved by reacting Hf and B at 1700 C and 95 MPa for 10 min in vacuo using SPS,59 instead of HP. For comparison, HfB2 ceramics produced from commercially available powder could achieved only 62% and 87% densities at 1800 C with 30-85 MPa for 10 min, and at 1900 C with 80-95 MPa for 10 min,59 respectively. The reaction between Hf and B occurred at 1100 C, while the completion of the reaction extended over a relatively wide temperature range.59 However, the association of the reaction with densification was absent during the reactive sintering. Densification was observed only at a temperature where the conversion to the diboride was complete. In addition, Zhao et al.60 showed that the reactive sintering of Zr, Si and B4C precursors could be conducted by SPS. The reactive SPS required a lower temperature of ≥1450 C with a shorter holding time of 3 min. The resulting composites had a finer and more homogeneous microstructure, compared with that from RHP. Thus, the              simultaneous synthesis and consolidation of the Zr, B, or B4C and Si precursor powders, i.e. reactive sintering, could produce the densified ZrB2 -based composites at a lower temperature by using either HP or SPS, as compared with direct consolidation of commercially available powders.  2.4. Pressureless sintering (PS)  In studies that were conducted in the 1970s and earlier, densification of ZrB2 ceramics was only accomplished by HP.5,24,26,27 Because of the extreme pressures required for densification, pressureless sintering of ZrB2 was considered unlikely or impossible until the late 1980s, when studies of pressureless sintering actually began to show results. Compared with HP, the development of a PS process would enable almost-net-shape processing of ceramic parts with complex geometries using standard powder-processing methods, thus reducing processing costs. Various additives have been used to improve densification of ZrB2 . In general, the additives used can be divided into main two groups: liquid phase formers, and reactive agents. Table 4 summarizes the PS conditions, agents used, grain size and final density of the resulting ZrB2 ceramics. Liquid phase formers include refractory metals, such as Ni, Fe, Co, and Mo,61-63 as well as disilicides of transition metals, 64,65 and ZrSi2 .20 Cech et al.61 used Ni, Co, Fe such as MoSi2 and Re to produce an almost fully densified ZrB2 at 2000 C that addition of ≥2 wt% of metals was required to bring about and 2200 C in vacuo or in an argon atmosphere. They found adequate sintering, because formation and continuous action of a liquid phase occurred only at higher contents of added metals. The additions are more efficient for producing adequate sintering in an argon atmosphere than in vacuo because of loss of the added metals in the vacuum from volatilization. Lattice parameter measurements showed a gradual decrease in the crystal lattice dimensions during sintering, resulting from substitution of zirconium atoms in the ZrB2 lattice by the atoms of the added metals. Obviously, the mode of action of these metallic additives that influence sintering is associated with an appreciable contraction of the ZrB2 crystal lattice. Presumably, the contraction of the ZrB2 crystal lattice affected the surface free energy, and, consequently increased the driving force for densification. A study63 in TiB2 with Ni, NiB, and Fe showed that a relative density exceeding 94% was obtained at a temperature ≥1500 C           \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  1001  Table 4  Pressureless sintering conditions, grain size, and final densities of the pressureless sintered ZrB2 ceramics with various additives.  Compositions  Remarks  Pressureless sintering conditions  Final density (%)  Grain size (\\u242em)  References  ZrB2 -20 vol% MoSi2 ZrB2 -4 wt% MoSi2 ZrB2 -10 vol% ZrSi2 ZrB2 -20-40 vol% ZrSi2 ZrB2 -4 wt% B4C ZrB2 -4 wt% B4C ZrB2 -4 wt% WC ZrB2 -2 vol% WC ZrB2 -1.7 wt% C ZrB2 -4 wt% B4C ZrB2 -4 wt% B4C ZrB2 -4 wt% B4C-0.5 wt% C ZrB2 -2 wt% B4C-0.5 wt% C  Ball-milled  1850                             C/30 min/Ar  99.1  2-3  65  Ball-milled  2250  C/120 min/Ar  97.7  3-11  66  Ball-milled  1650  C/60 min/Ar  95.7  -  20  Ball-milled  1650  C/60 min/Ar  >99  -  20  Attrition-milled  1850  C/60 min/vacuum  >98 97 95  12  67  None  2050  C/240 min/vacuum  5-6  67  Attrition-milled  2050  C/240 min/vacuum  -  67  Attrition-milled  2150  C/540 min/He  98  9  71  Attrition-milled  1900  C/120 min/Ar  >99  14  70  None  2050  C/120 min/Ar  94  6  69  Attrition-milled  1850  C/120 min/Ar  100  8  69  Attrition-milled  1850  C/120 min/Ar  >99  <4  69  None  1900  C/120 min/Ar  100  4  69  without external pressure, but was accompanied by significantly exaggerated grain growth. The addition of carbon inhibited grain growth, but also significantly increased the porosity. The authors hypothesized that the densification process occurs by redistribution followed by dissolution-reprecipitation in the nickel-rich melt, and that the grain growth was caused by surface diffusion in a titanium-oxide-rich surface layer.63 In addition, Kislui and Kuzenkova62 found that with Mo additions up to 15 wt% the energy of activation of the densification process drops below 380 kJ/mol from 680 kJ/mol because Mo was incorporated into a ZrB2 solid solution. As a result, the addition of Mo activates diffusion processes during sintering, effectively promoting densification. This densification behaviour of Mo is also found in ZrB2-20 wt% SiC mixture powder that could be sintered without pressure to a relative density exceeding 97.7% at 2250 C for 120 min when 4 wt% Mo was added.66 Other studies found that the addition of 20 vol% MoSi2 produced the almost fully densified ZrB2 at 1850 C for 30 min without external pressure.64,65 the additions of ≥20 vol% ZrSi2 further reduced Furthermore, the densification temperature; the full density was obtained at 1650 C for 60 min.20 Differing from the previously mentioned liquid phase formers, reactive agents act as densification aids by reacting with the oxide impurities present on the surface of starting particles (such as ZrO2 and B2O3 ) which inhibit densification. The main reactive agents used recently include B4C,67-69 C,68-70 and/or WC.67,71 It was found that the addition of 4 wt% B4C produces almost the fully dense ZrB2 compact at 1850 C for 60 min in vacuo for attrition-milled powder, without externally applied be sintered to a relative density of 95% at 2050 pressure.67 In contrast, ZrB2 containing only 4 wt% WC could C for 240 min. (2 vol%) allows sintering ZrB2 powder to an almost fully denChamberlain et al.71 also showed that the introduction of WC sified state at 2150 C for 180 min. They showed that elimination of oxide impurities on ZrB2 particles surface by the reactions of B4C or WC with ZrO2 was the key to densification. The at a temperature ≥1200 above-mentioned reactions are thermodynamically favourable C for B4C, but >1500 C for WC. As a result, the discrepancy in densification temperature between the two agent-doped ZrB2 ceramics is likely associated with their different onset temperatures for the reactions.                          Moreover, grain size measurement showed that excess B4C restricts the grain growth during sintering. A similar densification and grain growth inhibition effect of B4C was also reported in ZrB2 -containing B4C and carbon, either alone or in combination.69 In addition, the densification effect of B4C depended on the starting ZrB2 powder size.69 ZrB2 with a particle size of 2 \\u242em allows sintering to a density of 95% at 2050 C for 120 min. For comparison, ZrB2 could be achieved size was reduced to 0.5 \\u242em by an attrition milling. Furtherwith full densification at 1850 C for 60 min when the particle more, the densification is more effective for a combination of B4C and C than for B4C alone. Using a combination of B4C and C, the same ZrB2 powder (2 \\u242em) could be sintered to almost full density at 1900 C for 120 min. However, the additional densification effect of carbon does not appear in the reduced particle size ZrB2 powder (0.5 \\u242em), which could be sintered to a full density at 1850 C for 60 min using either B4C or a combination of B4C and carbon. Recently, Zhu et al.70 coated a carbon layer surface of ZrB2 particles using a phenolic resin as the carbon source. They found that the fully densified ZrB2 compact could be sintered without pressure at 1900 C for 120 min, as the carbon content is more than 1.0 wt% in the coated ZrB2 powders. For comparison, only a relative density of 70% was obtained for the C-uncoated ZrB2 powder under the same PS condition.                 3. Mechanical behaviours  3.1. Young’s modulus  Table 5 summarizes the Young’s modulus of ZrB2 ceramics with and without sintering additives. The Young’s modulus ranges from 350 GPa to 530 GPa, depending on porosity and additives; for a fully densified ZrB2 without additive, it is equal to 498 GPa.11 Historic studies have shown the Young’s modulus of the fully densified hot-pressed polycrystalline ZrB2 to to 500 GPa.72,73 The additions of Ni, AlN, Si3N4 , be equal B4C and C affect the Young’s modulus of ZrB2 ceramics; for is 496 GPa,29 higher ZrB2-4 wt% Ni, it than that of ZrB2 of the same density.71 The Young’s modulus of fully densified solids is determined principally by interatomic forces, which decrease sharply with the interatomic distance.74 The addition of  \\x0c', '1002  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  Table 5  Young’s modulus, fracture toughness, and 4-point flexural strength of the ZrB2 ceramics with and without additives.  Compositions  Process  Grain diameter (\\u242em)  Relative density  (%)  Young’s modulus  (GPa)  Fracture toughness (MPa m1/2 )  Flexural strength  (MPa)  References  ZrB2 ZrB2 ZrB2 ZrB2 ZrB2 -4 wt% Ni ZrB2 -4.6vol%AlN ZrB2 -5 vol% Si3N4 ZrB2 -4 wt% B4C ZrB2 -4 wt% B4C ZrB2 -2 wt% B4C-1 wt%C  HP  7.7  87  346  2.4  350  32,33  HP  6.0  99.8  489  3.5  565  11  HP  6.8  91  417  4.8  457  21  PS  9  98  454  -  444  71  HP  5-15  98  496  2.8  371  29  HP  4.2  92  407  3.1  580  33  HP  3  98  419  3.8  419  32  PS  6  94  500  3.3  489  69  PS  8  100  530  3.1  370  69  PS  4.1  >99  507  3.5  473  68,69  Ni caused a decrease in the crystal lattice dimensions of ZrB2 , as a result of substitution of Zr atoms in the ZrB2 lattice by atoms of the added Ni.61 Therefore, the increase of Young’s modulus due to Ni addition is associated with the decrease in the crystal lattice dimensions of ZrB2 . B4C alone or a combination of B4C with C also led to an increase in Young’s modulus. In contrast, AlN and Si3N4 additions led to a lower Young’s modulus. The different changes with the additives are associated with the different grain-boundary phase developed between ZrB2 grains,32,33,69 which result from the interactions of the additives with impurities on the ZrB2 particles surfaces, because the grain-boundary phase affects the Young’s modulus of ceramics.34,75 Fig. 2 is a plot of Young’s modulus as a function of the volume fraction of added of SiC, MoSi2 and ZrSi2 to the ZrB2 -based composites. The Young’s modulus of the three compositions decreased with amount added, in the order ZrB2 -SiC > ZrB2 -MoSi2 > ZrB2 -ZrSi2 , the maximum modulus value of SiC (475 GPa76 ), middle resulting from (440 GPa77 ), and minimum value of ZrSi2 (235 GPa40 ). For a fully densified composite, Young’s moduvalue of MoSi2 lus, Ec , may be described by the rule of mixtures78  Ec =  n(cid:2)  i=1  EiVi  (2)  Fig. 2. Young’s moduli as a function of additive content  for  the hot-pressed  ZrB2 -based composites with SiC, MoSi2 and ZrSi2 additives.  where Ei is Young’s modulus of ith constituent phase, and Vi is volume fraction of ith constituent phase, and n is total number of constituent phases. With E1 = 500 GPa (ZrB2 ),73 E2 = 475 GPa (SiC),76 E3 = 440 GPa (MoSi2 ),77 and E4 = 235 GPa (ZrSi2 ),40 Young’s moduli predicted by Eq. (2) are also drawn with three different lines in Fig. 2. A comparison between the measured and predicted values found that Young’s modulus of pore-free ZrB2 -based composites obeys the rule of mixtures. For the ZrB2-10 vol% ZrSi2 composition, the predicted Young’s modulus is higher than that measured experimentally, as a result of the presence of pores (RD: 96.6%). Fig. 3 shows the effect of porosity on Young’s modulus measured in ZrB2 -SiC and ZrB2-ZrC-SiC compositions. It has been found that Young’s modulus of the pores-containing ZrB2 -based composites is mostly dominated by the porosity. For ceramic materials, the use of a linear empirical dependence has been recommended. Assuming that the effect of pore structure and shape on Young’s modulus is neglected, Young’s modulus, E, can be given by79 E = E0 (1 − αP )  (3)  where E0 is Young’s modulus of pore-free materials, α is a constant, and P is the volume fraction of porosity in the material.  Fig.  3. Young’s moduli  as  a  function  of  porosity  for  the  hot-pressed  ZrB2 -30 vol% SiC and spark plasma sintered ZrB2 -ZrC-SiC composites.  \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  1003  The E0 and α values are obtained from the E-P plots. In addition, the effect of pore structure and shape on Young’s modulus have been investigated by other researchers.80,81 These studies demonstrated that the local elastic moduli decrease when the pore shape changed from spherical to oblate, as a result of the increased stress concentration around the pore. Therefore, the effective elastic modulus measured is lower for the case of nonspherical pores than for the case of spherical pores. Also, it has been found that elastic properties have a different sensitivity to porosity, regardless of pore shape. Bulk modulus is most affected by porosity; Young’s modulus is the next most-affected modulus, followed by the shear modulus (which is slightly less sensitive than Young’s modulus). In contrast, Poisson’s ratio is insensitive to additives as well as to porosity, and it remains almost the constant for the various ZrB2 -based composites.19,21,34,50  3.2. Fracture toughness  Fracture toughness of the ZrB2 ceramics with and without additives is also summarized in Table 5. Fracture toughness was in the range 2.4-4.8 MPa m1/2 . For the ZrB2 ceramics with, low fracture toughness, intragranular cracking is common with the cracks propagated across ZrB2 grains without being deflected along the grain boundaries. For the ZrB2 with high fracture toughness, intergranular cracking is partially present. In particular, for the case of the fracture toughness of 4.8 MPa m1/2 ,21 the intergranular cracking was the dominant crack propagation type, with the crack propagated along the grain boundaries (Fig. 4(a)). Thus, large grains and intergranular cracking are required for increasing fracture toughness. Also, the fracture toughness measured is larger in the ZrB2 with nitrides and/or carbides than with metallic additives. This difference seems to be associated with a larger tensile residual thermal stress in the ZrB2 ceramics with metallic additives than in the ZrB2 ceramics with carbide and/or nitride ceramic additives. Table 6 summarizes the fracture toughness of the ZrB2 -based composites with SiC, MoSi2 and ZrSi2 additives. Rezaie et al.15 found that the fracture toughness of ZrB2 -based composites with SiC is dominated by the SiC particles size and distribution in the composites. Increased fracture toughness produced by SiC addition is attributed to the crack deflection that occurs near the SiC particles and/or at ZrB2 /SiC interfaces. The interactions of the crack with the microstructure are most likely controlled by the complex residual stress state that develops during cooling from the processing temperature; that in turn is caused by the thermal expansion mismatch between the ZrB2 and SiC particles. The contribution of crack deflection to increasing fracture toughness depends on the total number of crack deflections and the crack deflection angle, i.e. crack propagation path. Crack deflection is enhanced in larger diameter grains. The increase of fracture toughness that accompanies crack deflection is also associated with the elastic and/or frictional bridging mechanism of grains. Elastic bridging is enhanced in larger diameter grains, while the frictional bridging mechanism is in operation only when crack deflection and grain pullout occur, which is prevalent in smaller diameter grains. Fig. 5 is a plot of fracture toughness as a function of the ratio of ZrB2 to SiC grain size for ZrB2-30 vol%  Fig. 4. Typical  cracking behaviours  intergranular cracking behaviour,  for hot-pressed (a) ZrB2 (b) ZrB2 -20 vol% MoSi2 -5 vol% SiC, and (c) ZrB2 -20 vol% MoSi2 -20 vol% SiC, showing crack deflection at small SiC particles and break of large SiC particles.  ceramic with  SiC composites. Note that these data are taken in the same test environments in order to avoid the effect of changing test con ditions. It was found that fracture toughness increased with the ratio of ZrB2 to SiC grain size. This increase of fracture toughness is attributed to the two major causes. One is large ZrB2 grains enhanced elastic and/or frictional bridging contributions to increasing toughness. Another is small SiC grains increased the number of crack deflections and pullout grains. Therefore, large ZrB2 grains and/or small SiC grains are required to improve  \\x0c', '1004  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  Table 6  Young’s modulus, Fracture toughness, 4-point flexural strength of the ZrB2 -based composites with SiC, MoSi2 and ZrSi2 additives.  Compositions (vol%)  Relative density (%)  Measured method  Fracture toughness (MPa m1/2 )  Young’s modulus  (GPa)  Flexural  strength (MPa)  References  ZrB2 -10SiC ZrB2 -20SiC ZrB2 -30SiC ZrB2 -30SiC ZrB2 -30SiC ZrB2 -30SiC ZrB2 -10MoSi2 ZrB2 -20MoSi2 ZrB2 -30MoSi2 ZrB2 -40MoSi2 ZrB2 -20MoSi2 -5SiC ZrB2 -20MoSi2 -10SiC ZrB2 -20MoSi2 -20SiC ZrB2 -20MoSi2 ZrB2 -10ZrSi2 ZrB2 -20ZrSi2 ZrB2 -30ZrSi2 ZrB2 -40ZrSi2  93.2  Four-point bending  4.1  450  713  11  99.7  Four-point bending  4.4  466  1003  11  99.4  Four-point bending  5.3  484  1089  11  99.8  Four-point bending  4.6  520  909  14  97.2  Four-point bending  5.5  516  1063  15  99.5  Four-point bending  4.5  505  804  15  99.7  Indentation technique  3.7  490  799  19  99.8  Indentation technique  2.8  472  749  19  99.8  Indentation technique  2.6  473  756  19  99.7  Indentation technique  3.1  448  790  19  100  Indentation technique  3.4  476  862  19  97.3  Indentation technique  3.4  465  554  19  94.6  Indentation technique  3.4  461  368  19  99.1  Indentation technique  2.3  489  531  65  96.6  Indentation technique  3.8  432  483  21  99.1  Indentation technique  4.4  445  556  21  99.8  Indentation technique  4.4  427  555  21  99.2  Indentation technique  3.9  397  382  21  the toughness of ZrB2 -based ceramics. Recently, a toughness of 6 MPa m1/2 was reached in the ZrB2-ZrC-SiC system by optimizing the combination of composition with microstructure.48 For the ZrB2-MoSi2 system, fracture toughness is in the range of 2.3-3.7 MPa m1/2 and decreased with the content of MoSi2 .19 Differing from the ZrB2-SiC system, for the ZrB2 -MoSi2 system the crack propagated along ZrB2 phase boundaries and across the MoSi2 phase,19 thereby decreasing the toughness of these composites. For comparison, in the ZrB2-ZrSi2 system the range of fracture toughness values was 3.8-4.4 MPa m1/2 , tending to improve with ZrSi2 content.21 This discrepancy between the composite systems seems to be associated with larger ZrB2 grain size and smaller ZrSi2 grain size in the ZrB2 -ZrSi2 system as compared to the ZrB2-MoSi2 system.19,21 Furthermore, it is found that an addition of a small amount of SiC (5 vol%) led to an increase in fracture toughness of the ZrB2-MoSi2 composites, but the toughness  Fig. 5. Plots of fracture toughness as a function of the ratio of ZrB2 to SiC grain size for the ZrB-based composites with SiC additives.  remained almost the constant with further increasing SiC content (≥10 vol%).19 The crack deflection at the SiC/ZrB2 interface and the multiple cracking at crack tips contributed to an increase of toughness (Fig. 4(b)). Even with increasing SiC content, crack deflection still occurred only at the smaller individual SiC particles; the crack was across the larger and/or agglomerated SiC particles (Fig. 4(c)). Obviously, the fractured larger SiC particles were not effective in contributing to increase of toughness. The constant toughness with SiC content suggests that most of the added SiC particles were fractured during cracking at high SiC content (≥10 vol%). Thus, a more uniform dispersion of ultrafine SiC particles in the ZrB2 matrix is important for optimizing fracture toughness.  3.3. Flexural strength  Room-temperature 4-point flexural strengths of the ZrB2 ceramics with and without additives are also summarized in Table 5. It was found that the flexural strengths range from 350 MPa to 580 MPa, depending on grain size, additives, and on the relative density as well. The hot-pressed ZrB2 ceramic with AlN showed the highest flexural strength of 580 MPa, although the flexural strength of ZrB2 ceramic with Si3N4 , with 3 \\u242em only a relative density of 92% was obtained.33 For comparison, grain size and 98% density, was 419 MPa. This discrepancy is probably associated with an internal tensile stress at the grain boundaries upon cooling from the pressing temperature, which is in turn is caused by the thermal expansion mismatch between the ZrB2 and grain boundary phase or shrinkage of the intergranular amorphous phase.82 It is evident that the ZrB2 ceramic with Ni has the lowest room-temperature flexural strength, resulting from the extremely large thermal expansion mismatch between the ZrB2 and Ni.29 Fig. 6 is a plot of room-temperature 4-point flexural strength as a function of additive content for hot-pressed ZrB2 -based  \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  1005  Fig. 6. Plots of  room-temperature 4-point flexural  strength as a function of  additive content  for  and ZrSi2 additives.  the hot-pressed ZrB2 -based composites with SiC, MoSi2  Fig. 7. Plots of room-temperature 4-point flexural strength as a function of grain  size of secondary phase for the hot-pressed ZrB2 -based composites with SiC and MoSi2 additives.  composites with SiC, MoSi2 and ZrSi2 additives. Chamberlain et al.11 showed that the addition of 10, 20 and 30 vol% SiC led to the high room-temperature strengths of ZrB2 , cally 1000 MPa, with fine ZrB2 grain microstructure (average typigrain size: 2-3 \\u242em). Guo et al.19 found that the addition of 10-40 vol% MoSi2 increased room-temperature flexural strength of ZrB2 ceramics to over 700 MPa. Bellosi et al.18 reported similar results in the 15 vol% MoSi2 -containing ZrB2 ceramics produced by SPS and HP. They found that the flexural strength for MoSi2 -containing ZrB2 composites consolidated both by SPS and HP increased to 640 MPa and 700 MPa, respectively, corresponding to an average grain size of 1.4 \\u242em and 1.8 \\u242em. Comparing both the strength values measured in the ZrB2-SiC and ZrB2-MoSi2 , it appeared that the strength is higher in the ZrB2-SiC composites than in the ZrB2 -MoSi2 composites although the grain size of ZrB2 is larger in the former than in the latter. One exception was a lower strength for a 10 vol% SiC-containing ZrB2 composite, as a result of the presence of more pores (RD: 92%). On the other hand, the addition of ZrSi2 improved densification of ZrB2 ceramics and inhibited (average grain size: 2-3 \\u242em), but grain growth as well the flexural strength did not increase significantly. Conversely, the (30%). Hence, although ZrB2 grain size is an important factor addition of 40 vol% ZrSi2 led to a degradation of the strength affecting the room-temperature flexural strength of ZrB2 -based composites with carbide and disilicides, it is not the limiting factor. A similar conclusion for ZrB2 -based composites with SiC was also reported by Fahrenholtz et al.5 Recently, the effects of microstructure and SiC grain size on room-temperature flexural strength were examined in hotpressed 30 vol% SiC-containing ZrB2 ceramics by Zhu et al.14 and Rezaie et al.,15 both studies showed, based on the linear elastic fracture mechanics, that the critical flaw size correlates strongly with SiC particle size. Both studies concluded that the maximum SiC grain size in the ZrB2 -SiC composites is the strength-limiting factor, and that the strength is not correlated with average ZrB2 grain size. Similar to ZrB2-SiC composites,  the strong relationship of strength to maximum MoSi2 grain size was also found in ZrB2 -MoSi2 composites (Fig. 7). Thus, it was assumed that the maximum MoSi2 or ZrSi2 grain size dominated the room-temperature strength for ZrB2 -based composites containing MoSi2 or ZrSi2 additives.19,20 Addition of 5 vol% SiC to ZrB2 -MoSi2 can lead to a further increase of flexural strength.19 The addition of nano-sized SiC also improved the flexural strength of the ZrB2-SiC composites before and after oxidation for 10 h at 1400 C in air (Fig. 8).25 For comparison, after the oxidation, the strength of 20 vol% micron-sized SiC-ZrB2 decreased from 530 MPa before the oxidation to 500 MPa (Fig. 9).25 In particular, for single-phase ZrB2 ceramloss of strength reached 70% after oxidation for 10 h at ics, 1400 C in air. For the nano-sized SiC-ZrB2 composites, the increase of strength after oxidation was attributed to the presence of a thin, dense oxide layer on the surface of the post-oxidized samples (Fig. 10(a)). The presence of a thin oxide layer can heal the surface flaws without creating new cracks and defects at the        Fig. 8. Effect of thermal exposure at 1400     C for 10 h in air on room-temperature  4-point flexural strength for particles.25  the ZrB2 -based composites with nano-sized SiC  \\x0c', '1006  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  Fig. 9. A comparison of flexural strengths of the single-phase ZrB2 ceramics, micronand nano-sized SiC-ZrB2 composites before and after oxidation at 1400 C for 10 h in air.25     oxidized surface, and plays a significant role in strength improvement after exposure to an oxidation environment.83 In contrast, for the micron-sized SiC-ZrB2 composites, loss of strength due to exposure to high temperatures is attributed to formation of a  Fig. 11. Plots of 4-point flexural strength as a function of temperature for various  the ZrB2 -based composites with various additives.     thicker glassy layer and a defect at the interface between the oxide layer and bulk ZrB2-SiC composites (indicated by an arrow in Fig. 10(b)). Although additives, such as Ni, SiC, and nitrides and disilicides, could reduce the densification temperature of ZrB2 , the strength of the resulting composites is also degraded at hightemperature (Fig. 11), as a result of softening of the intergranular amorphous phase. For the ZrB2 -based ceramics with Ni and Si3N4 , the strength degraded significantly above 800 For the ZrB2 -based ceramics with the Ni additive, in particular, the strength decreased sharply and dropped almost to zero at 1200 C. For comparison, the ZrB2 -based composites with MoSi2 retained the constant strength at temperatures approaching 1200 C.18,65 However, the strength degraded rapidly above 1200 C, and the degradation was more rapid than in the ZrB2 ceramics without additive. One exception was the HfB2-SiC composite produced by SPS, which retained room-temperature strength up to 1500 C.18 Although the cause is not completely understood, it is assumed to be closely linked with the minimization of impurities during SPS. Thus, high-temperature strength may be improved by increasing the refractory index of the intergranular phase and minimizing impurities, as well as promoting crystallization of the intergranular amorphous phase.82  C.22,30              4. Physical behaviours  Single-phase ZrB2 and ZrB2 -based composites have high thermal and electrical conductivities among the transition metal carbides, nitrides and diborides. High electrical conductivity allows the fabrication of complex shapes using electrical discharge machining. In addition, high thermal conductivity could improve thermal shock resistance by reducing temperature gradients and thermal stress within the materials. However, the thermophysical and electrical properties of ZrB2 -based composites have not been extensively investigated. This section focuses on recent studies in thermal conductivities and electrical conductivities of ZrB2 ceramics and ZrB2 -based composites.  Fig. 10. Typical  the hot-pressed ZrB2 -based composites with nanoand micron-sized SiC particles after oxidation exposure at 1400 C  fracture surfaces of     for 10 h in air: (a) 20 vol% nano-sized SiC-ZrB2 , and (b) 20 vol% micron-sized SiC-ZrB2 composite.25  \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  1007  Table 7  Thermal conductivity and electrical conductivity of the ZrB2 ceramics and ZrB2 -based composites with various additives. −1 K −1 ) Electrical conductivity σ (×104 , \\x01−1 cm −1 ) Thermal conductivity k (W m  Composition (vol%)  10.3  4.5  6.7  4.2  4.2  13.5  6.2  14.3  6.7  1.61  4.52  0.92  3.00  8.11  7.22  7.48  7.35  6.04  4.07  10.7  11.6  11.9  9.6  References  8  84  32,33  84  33  29  29  32  32  48  48  48  48  87  87  87  87  87  87  20  20  20  20  ZrB2 ZrB2 ZrB2 ZrB2 -30% SiC ZrB2 -4.6AlN ZrB2 -4Ni (wt%) ZrB2 -13B4C-4Ni (wt%) ZrB2 -5Si3N4 ZrB2 -20SiC-4Si3N4 ZrB2 -33.3ZrC-33.3SiC (mol.%) ZrB2 -15ZrC-15SiC (mol.%) ZrB2 -55ZrC-30SiC (mol.%) ZrB2 -15ZrC-30SiC (mol.%) ZrB2 -10MoSi2 ZrB2 -20MoSi2 ZrB2 -40MoSi2 ZrB2 -40MoSi2 -5SiC ZrB2 -40MoSi2 -10SiC ZrB2 -40MoSi2 -20SiC ZrB2 -10ZrSi2 ZrB2 -20ZrSi2 ZrB2 -30ZrSi2 ZrB2 -40ZrSi2  4.1. Thermal conductivity  58.2  56.4  -  62.1  -  -  -  -  -  72.6  85.6  51.8  89.0  87.6  82.8  76.1  80.8  84.6  88.9  98.3  96.8  86.8  74.2  −1  −1  −1  −1  Table 7 summarizes the thermal conductivities of the ZrB2 ceramics and ZrB2 -based composites with various additives. The thermal conductivity was generally in the range of 50 W (m K) and 60 W (m K) for the ZrB2 ceramics.8,84 Also, an early study85 of a polycrystalline ZrB2 showed that thermal conductivity of the ZrB2 ceramics could achieve a higher value of 84 W (m K) −1 . For single-crystal ZrB2 , the thermal conductivity value was measured to be 140 W (m K) in the basal direction and 100 W (m K) along the c-axis.86 Thus, the thermal conductivity of the ZrB2 -based composites should −1 and 140 W (m K) −1 . Recently, in the be between 50 W (m K) thermal conductivity measurements in various ZrB2-ZrC-SiC composites consolidated by SPS, it appeared that the conduc−1 , tivity was in the range of 38 W (m K) and 93 W (m K) showing a strong compositional dependence.48 The thermal conductivity of the ZrB2-ZrC-SiC composites increased with increasing ZrB2 as well as SiC content, whereas the conductivities decreased with increasing ZrC content. The compositional dependence of thermal conductivity is also observed in ZrB2 -ZrSi2 and ZrB2-MoSi2 -SiC composite materials.20,87 It is found that the increase of MoSi2 or ZrSi2 led to a decrease of thermal conductivity of the composites. Conversely, increasing SiC enhanced thermal transport in the ZrB2-MoSi2-SiC materials; therefore, high thermal conductivity, but the content of added SiC must exceed 5 vol%. It is known that the thermal conductivity of the composites depend on the thermal conductivity of the components and the interfacial thermal resistance between the components. In the ZrB2 -based composites, the addition of the second phase with higher thermal conductivity than the ZrB2 phase, such as  −1  SiC, led to decreased resistance for the heat flow through the components and their interfaces, thereby increasing thermal conduction. In particular, in the case of the composition with high concentration with SiC, the phase with higher conductivity could form a network-like structure which improved heat capacity, and enhanced heat transport as well, resulting in high thermal conductivity. For comparison, the addition of the second phase with lower thermal conductivity, such as ZrC, MoSi2 and ZrSi2 , led to a lower thermal conductivity, probably resulting from the high thermal transport resistance of the second phase, as well as from a higher intergranular thermal resistance. Fig. 12 is a plot of thermal conductivity as a function of test temperature for the ZrB2 ceramics and ZrB2 -based composites  Fig. 12. Plots of thermal conductivity as a function of test  temperature for the  single-phase ZrB2 ceramics and ZrB2 -based composites with SiC additives.  \\x0c', '1008  S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011     with SiC reported in the studies.84,88 Thermal conductivity was calculated using measured thermal diffusivity determined by the laser flash method, bulk density, and heat capacity. It was found that the thermal conductivity of the ZrB2 ceramics showed a considerable weak temperature dependence; in particular, the conductivity was almost the constant in the range of 400 C to 1700 C. In contrast, for the ZrB2 -based composites with SiC, thermal conductivity showed strong temperature dependence. It was found that the thermal conductivity decreased with temperature, regardless of SiC content. Zimmenmann et al.84 found that the ZrB2 grain size affects the temperature dependence of the thermal conductivity, and that the conductivity decreased with reduction in grain size. They suggested that the ZrB2 ceramics and ZrB2 -based composites for use under thermal loading conditions should be examined to determine the optimal grain size to balance the need for thermal transport and strength.     Fig. 13. Plots of electrical resistance as a function of content of TiB2 ZrB2 -TiB2 composites at various temperatures.89  for the  −1  −1 for ZrB2 ,84,85,89,90  range of 1.0 × 10 −3 K −1 to 6.4 × 10 −3 K and the TCR is approximately 1.4 × 10 −3 K −1 for a solid solution of (Zr1−xTix )B2 (0 ≤ x ≤ 1).89 The addition of SiC increased the TCR of ZrB2 ceramics. Tye and Clougherty85 showed that 4.8 × 10 −3 K the TCR value was for ZrB2-20 vol% SiC composites, whereas Zimmenmann et al.84 reported that the measured TCR value was 2.5 × 10 −3 K −1 for ZrB2 -30 vol% SiC. Also, Jimbou et al.91,93 and Takahashi et al.92 found that an electrical percolation threshold is observed in the ZrB2 -SiC system. The electrical resistivity of ZrB2 -SiC composites sharply increased with SiC content over 70 vol%, as shown in Fig. 15. Presumably, below the critical content of SiC, the network-like structure formed by the ZrB2 phase with high-electrical conductivity provides an electrical path with lower resistance, thereby retaining high electrical conductivity. The electrical resistivity values measured in the ZrB2-SiC composition (SiC content: <70 vol%) have below an order of 6 in magnitude, compared with that of pure SiC. Furthermore, the grain diameter and  4.2. Electrical conductivity  The ZrB2 ceramics and ZrB2 -based composites are electrical conductors, which exhibited metallic-like electrical conductivity (Table 7). The electrical conductivity of the pure ZrB2 ceramics without additives is in the range of 4.5 × 104-10.3 × 104 (\\x01 cm) −1 .8,32,33,89 The discrepancy in the data reported by different authors is probably due to different impurity levels, processing route, and measurement methods. For the ZrB2 ceramic with additives, the addition of Ni decreased electrical resistance, thereby enhancing electrical conductivity. For comparison, the additions of carbides and nitrides increased electrical resistance, thereby reducing electrical conductivity.29,33 One exception was an electrical conductivity of the ZrB2 -based ceramics with Si3N4 comparable to that of the pure ZrB2 ceramics although Si3N4 is an insulator.32 In addition, the electrical conductivity of ZrB2 depended on composition. Rahman et al.89 showed that the electrical conductivity in the ZrB2-TiB2 system decreased with increasing TiB2 , where the zirconium atoms in the ZrB2 lattice were substituted by the titanium atoms to form a solid solution of (Zr1−xTix )B2 (0 ≤ x ≤ 1). As a result, the lower electrical conductivity is due to the higher electrical resistivity of the TiB2 phase compared the ZrB2 phase.89 to that of In addition, they found that the compositional dependence appears to follow linear behaviour (below 200 at low temperatures C, Fig. 13), but becomes nonlinear at high temperatures. Furthermore, the temperature dependence of the electrical resistivity of the ZrB2 ceramics and ZrB2 -based composites are similar to that of metals; the resistivity increased linearly with temperature, examples of which are shown in Fig. 14. However, the addition of SiC led to an increase of electrical resistivity of ZrB2 ceramics; the increase was enlarged significantly with increasing temperature. The thermal coefficient of electroresistance (TCR), β, was given by90     β =  1  dρ(T )  ρ298  dT  (4)  where ρ298 is the room temperature electrical resistivity, and T is the absolute temperature. The TCR values were found to be in the  Fig. 14. Plots of electrical resistivity as a function of temperature for the ZrB2 ceramics and ZrB2 -30 vol% SiC composite.  \\x0c', 'S.-Q. Guo / Journal of the European Ceramic Society 29 (2009) 995-1011  1009  at lower temperatures, compared with HP. Pressureless sintering of ZrB2 -based composites was possible by addition of B4C, C, MoSi2 and ZrSi2 . Spark plasma sintering (SPS) could densify various compositions of ZrB2 -based composites at a lower temperature and very short soaking time, compared to the HP, RHP and PS methods. (2) Young’s modulus of the ZrB2 -based ceramics was in the range of 300-500 GPa, strongly dependent on additives and porosity. Young’s modulus of porosity-free ZrB2 -based composites could be predicted by the rule of mixtures. (3) Fracture toughness of the ZrB2 -based ceramics was in the range of 2-6 MPa m1/2 . The fracture toughness is dominated by secondary phase particle size and distribution. Larger ZrB2 grain and small second phase grain could increase fracture toughness, as a result of crack deflection at the grain-boundary interfaces. This is associated with the complex residual stress within grains and at the grainboundaries. (4) Flexural strength of the ZrB2 -based ceramics was significantly increased by additions of second phase particles. The secondary phase particle size and distribution governed the room-temperature strength. The addition of fine SiC particles showed the highest room temperature strength. In addition, the addition of nano-sized SiC improved strength of ZrB2-SiC composites after oxidation of 10 h at 1400 C in air, while the strength of ZrB2 -based composites with micron-sized SiC decreased. Further improvement in strength may be achieved through a more uniform dispersion of nano-sized particles in ZrB2 matrix. (5) Thermal conductivity of the ZrB2 -based ceramics was in −1 , depending on composithe range of 30-140 W (m K) tion as well as grain size. The addition of SiC increased thermal conductivity, whereas the additions of MoSi2 and ZrSi2 decreased the conductivity. The thermal conductivity remained almost constant with temperature for the ZrB2 ceramics in the range of 400-1700 C, whereas the conductivity of ZrB2 -based composite with SiC decreased with temperature. (6) Electrical conductivity of the ZrB2 -based composites was in the range of 1 × 104 (\\x01 cm) to 12 × 104 (\\x01 cm) −1 , depending on composition. The additions of SiC and MoSi2 led to decrease in the electrical conductivity, while the addition of ZrSi2 (≤30 vol%) had not the noticeable effect on the electrical conductivity.  −1        References  1. Upadhya, K., Yang, J.-M. and Hoffmann, W. P., Materials for ultrahigh temperature structural applications. Am. Ceram. Soc. Bull., 1997, 76(12),  51-56.  2. Berton, B., Bacos, M. P., Demange, D. and Lahaye, J., High-temperature  oxidation of silicon carbide in simulated atmospheric re-entry conditions. J. Mater. Sci., 1992, 27, 3206-3210.  3. Scott, H. G., Phase relationships in the zirconia-yttria system. J. Mater. Sci., 1975, 10, 1527-1535.  4. Kajihara, K., Yoshizawa, Y. and Sakuma, T., The enhancement of superplas tic flow in tetragonal zirconia polycrystals with SiO2 -doping. Acta Metall. Mater., 1995, 43(3), 1235-1242.  Fig. 15. 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Takahashi, K. and Jimbou, R., Effect of uniformity on the electrical resistivity of SiC-ZrB2 ceramic composites. J. Am. Ceram. Soc., 1987, 70(12), C-369-C-373.  72. Wiley, D. E., Manning, W. R. and Hunter  Jr., O., Elastic properties of  93.  Jimbou, R., Suzuki, Y. and Takahashi, K., Effect of dispersion on resistivity  polycrystalline TiB2 , ZrB2 and HfB2 from room temperature to 1300 K. J. Less-Common Met., 1969, 18, 149-157.  of SiC-ZrB2 electroconductive ceramic composites. J. Ceram. Soc. Jpn., 1990, 98(2), 225-230.  \\x0c']"
},{
  "_id": 37,
  "PDF": "Densification, Mechanical Properties, and Oxidation Resistance of TaC–TaB2 Ceramics.pdf",
  "Text": "['Densification, Mechanical Properties, and Oxidation Resistance of  TaC-TaB2 Ceramics  Xiaohong Zhang,*,w  Gregory E. Hilmas,* and William G. Fahrenholtz**  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, Missouri  65409  The densification of tantalum carbide (TaC) was enhanced by adding 10 wt% (11 vol%) TaB2, reaching 98.6% relative density by hot pressing at 21001 or 22001C using a 30 MPa applied pressure. X-ray diffraction analysis identified two phases, TaC  and tantalum diboride (TaB2), with no peak shifts, solid solubility was not significant at these temperatures. Me indicating a  chanical properties were measured for TaC-10 wt% (11 vol%) TaB2 hot pressed at 21001C and compared with monolithic TaC hot pressed at 23001C. The Young’s modulus increased from 472 GPa for monolithic TaC to 543 GPa for TaC with 10 wt% (11  vol%) TaB2 additions. Vickers’ hardness increased from 14.1 to 19.4 GPa. Fracture toughness values were comparable, in the 3.4-3.5 MPa . m1/2. The was 600 MPa, comparable to the flexure strength of monolithic  range  of  average  flexure  strength  TaC  (686 MPa).  Thermal  gravimetric  analysis  indicated  that  the TaC-10 wt% (11 vol%) TaB2 oxidized at a slightly higher temperature than monolithic TaC. Together, the results  indicate that TaB2 is an effective sintering aid for the densification of TaC.  I.  Introduction  TANTALUM carbide (TaC) and tantalum diboride (TaB2) belong to the family of ultra-high-temperature ceramics (UHTCs). These materials have attracted interest in recent years  due to the increased demand for higher performance materials  for applications in extremely harsh environments such as pro pulsion systems for future hypersonic vehicles.  Monolithic TaC has been difficult to densify even by hot pressing at extremely high temperatures (24001C) because of its  highly covalent bonding, low self-diffusion coefficient, oxide impurities, and rapid grain growth at high temperatures.1 Liquid phase sintering,  through the addition of  transition metals, has  been shown to lower the densification temperature of TaC by 9001C for same density.2,3 However,  about  the  rapid grain  growth caused by the presence of a liquid phase  resulted in  the entrapment of porosity within the grains, which limited the  density that could be achieved. Further, the presence of metallic  phases at the grain boundaries is likely to degrade the high-temperature properties of the TaC ceramic.4 The formation of solid  solutions between TaC and other carbides (ZrC and HfC) has  been shown to improve densification by hot pressing at extreme temperatures (above 25001C).5 However,  the high temperature  required to achieve high density again resulted in grain growth.  Physical pinning of grain growth without forming a solid solu tion, such as through the addition of a UHTC diboride,  is an other approach that may enhance the densification of TaC. The  addition of refractory diborides may also provide better oxida tion resistance in addition to enhancing densification of TaC  because borides are more oxidation resistant than carbides, general.6 The study of system7 has  in  the TaC and TaB2 that TaC and TaB2 exhibited almost no solid solubility in each other up to 21001C. Further, the solubility of TaB2 in TaC is o3 wt% at 24001C and about 7 wt% at 27301C.7  shown  In the present paper, TaC containing 10 wt% (11 vol%) TaB2 was hot pressed to 498% relative density at 21001C. The me chanical properties of TaC-10 wt% (11 vol%) TaB2 were measured and compared with hot-pressed monolithic TaC having a  relative density of 94%. Oxidation resistance was analyzed by  thermal gravimetric analysis (TGA) for both materials and was  compared with TaC and TaB2.  II.  Experimental Procedure  TaC (TA-301, 1-5 mm, Atlantic Equipment Engineers, Bergenfield, NJ) and TaB2 (0.4 mm, 95% TaB2 and 5% TaC, synthesized at 15001C using Ta2O5, B4C, and graphite8) powders were mixed by attrition milling in hexane with SiC grinding media for  30 min. After drying by rotary evaporation, the powder mixture  was hot pressed (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 21001 or 22001C under a helium atmosphere. The  details of where.1  the hot-pressing schedule have been described else The bulk densities of  the hot-pressed billets were measured  using the Archimedes method with kerosene as the immersion  medium. The true density was calculated according to the rule of  mixtures based on the nominal TaC and TaB2 contents. The relative density was then determined by the ratio of bulk density  to the true density. Microstructural analysis was performed us ing scanning electron microscopy (SEM, Hitachi S-570, Tokyo,  Japan). The grain sizes were calculated by measuring an average  of 200 grains using image analysis software (Image J, U.S. Na tional  Institutes of Health, Bethesda, MD). X-ray diffraction  analysis  (XRD, XDS 2000, Scintag Inc., Cupertino, CA) was  performed to determine phase  compositions. Peaks were  in dexed using PDF cards 35-0801 for TaC and 75-0966 for TaB2. The relative amounts of the phases were quantified using Riet veld refinement software (RIQAS, Materials Data Inc., Liver more, CA). TGA (STA-409C, Netzsch, Selb, Germany) was performed on rectangular bars (1.5 mm \\x02 2.0 mm \\x02 10 mm) at a heating rate of 51C/min in flowing air up to 15001C. Young’s  modulus was measured  by  the  impulse  excitation method  (Model MK4-I Grindosonic, J.W. Lemmens, St. Louis, MO).  Microhardness was measured using Vickers’ indentation (Model  Duramin-5, Struers Inc., Westlake, OH) by applying a load of  0.5 kg (4.91 N). Four-point bending strength was measured in a  mechanical  load frame (Model 5881, Instron, Norwood, MA) (2.0 mm \\x02 1.5 mm \\x02 25 mm) according to  using type A bars  I. Talmy—contributing editor  This work was financially supported by the U. S. Army Space and Missile Defense  Command under contract DASG60-03-1-0011.  *Members, The American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: xiaoh@mst.edu  **Fellow, The American Ceramic Society.  Manuscript No. 25022. Received July 23, 2008; approved September 12, 2008.  Journal  J. Am. Ceram. Soc., 91 [12] 4129 - 4132 (2008)  DOI: 10.1111/j.1551-2916.2008.02780.x  r 2008 The American Ceramic Society  4129  \\x0c', '4130  Communications of the American Ceramic Society  Vol. 91, No. 12  around 5 wt% TaC. When the TaC content of the TaB2 is taken into account, the starting composition was predicted to be 90.5  wt% (89.2 vol%) TaC and 9.5 wt% (10.8 vol%) TaB2.  (2)  Densification and Microstructure  The bulk density of TaC-10 wt% (11 vol%) TaB2 hot pressed at 21001 and 22001C was 14.1 g/cm3. The true density was calculated to be 14.3 g/cm3 using a volumetric rule of mixtures and  the measured composition. Based on these values,  the relative  density was then calculated to be 98.6% for ceramics hot pressed at either 21001 or 22001C. For comparison, a previous  study  found that the relative densities of the same TaC powder hot pressed at 21001 and 22001C were 85% and 89% after hot pressing at 21001 and 22001C, respectively.1 The enhanced dens ification of TaC with the TaB2 addition is likely due to physical pinning of grain growth by the second phase. In addition to  grain pinning, the presence of B2O3 on the TaB2 particles’ surface may facilitate grain rearrangement due to the formation of  a liquid phase, which would enhance densification. A similar  grain growth-pinning effect has been observed in the ZrB2-SiC and HfB2-SiC systems.10 In these systems, dispersed SiC particles have been shown to inhibit the grain growth and enhance  densification of the matrix phase.  SEM images of TaC-10 wt% (11 vol%) TaB2 hot pressed at 21001 and 22001C (Fig. 2) showed that the microstructure was  composed of darker TaB2 grains (circled) uniformly distributed in a lighter TaC matrix. The average grain size of TaB2 was 1.1 specimens hot pressed at 21001 and 22001C, and 1.5 mm for  respectively. The size of the TaC grains could not be determined  because the grain boundaries of TaC could not be distinguished  in the polished cross  sections. Repeated attempts at  thermal  etching to reveal the grain boundaries were not successful.  (3) Mechanical Properties  The mechanical properties (Young’s modulus, hardness, flexure  strength, and fracture toughness) for TaC-10 wt% (11 vol%) TaB2, hot pressed at 21001C, are included in Table I. Also included in Table I, for comparison, are the mechanical property data for monolithic TaC hot pressed at 23001C (94% relative  density). The Young’s modulus  for TaC-10 wt% (11 vol%)  TaB2 was 543 GPa. This value was close to the value of 538 GPa calculated using a volumetric rule of mixtures for a dispersed phase composite containing 10.8 vol% (9.5 wt%) TaB2 (E 5 551 GPa)8 in a TaC matrix (E 5 537 GPa).11 Vickers’ hardness in creased from 14.1 GPa for monolithic TaC to 19.4 GPa for  ness of  attributed to the  TaC-10 wt% (11 vol%) TaB2, which was higher density achieved for the composite and the higher hardthe dispersed TaB2 phase (24.5 GPa)12 compared with vol%) TaB2 was determined to be B3.4 MPa \\x01 m1/2, nearly idenmonolithic TaC. The fracture toughness for TaC-10 wt% (11 tical to the fracture toughness of monolithic TaC (B3.5 MPa \\x01 m1/2) and slightly lower than that of monolithic TaB2 (B4.5 MPa \\x01 m1/2).8 The average strength of TaC-10 wt% (11  Fig. 1.  X-ray diffraction patterns of tantalum carbide (TaC) and tan talum diboride (TaB2) as well as TaC-10 wt% TaB2 after hot pressing at 21001 and 22001C.  ASTM Standard C1161. Fracture toughness was determined by the Indentation Strength in Bending method,9 which involves  measuring the strength after Vickers’ indentation. A load of 5 kg  was applied to produce radial-median cracks on type A flexure  bars before fracturing specimens  in four-point bending to de termine the fracture toughness.  III.  Results and Discussion  (1)  XRD Analysis  XRD patterns of TaC (hot pressed at 22001C), TaB2 (synthesized at 16001C, hot pressed at 21001C), and TaC-10 wt% (11 vol%) TaB2 hot pressed at 21001 and 22001C are included in Fig. 1 for comparison. No peak shifts were observed in the XRD  patterns for any of the TaC-10 wt% (11 vol%) TaB2 specimens, suggesting that TaB2 did not dissolve into TaC at the temperatures used for processing of these materials. This result is consistent with the observation of Ordan’yan et al.7 that TaC and TaB2 are virtually insoluble in each other up to 21001C. The insolubility of TaB2 in TaC may be due to the difference in the crystal structures. TaB2 crystallizes in the hexagonal structure (AlB2 type) while TaC forms as a cubic structure (B1). The relative amounts of TaC and TaB2 in the hot-pressed TaC-10 wt% (11 vol%) TaB2 samples, as determined by Rietveld refinement, were 91 wt% (89.8 vol%) TaC-9 wt% (10.2 vol%) TaB2. This was close to the nominal starting composition (90 wt% TaC-10  wt% TaB2) and even closer to the starting composition considering the fact that the TaB2 synthesized at 15001C contained  2100°C  2200°C  TaB2  TaC  TaB2  10 µm  TaC  10 µm  Fig. 2. Microstructures of tantalum carbide (TaC)-10 wt% tantalum diboride (TaB2) hot pressed at 21001C (left) and 22001C (right).    \\x0c', 'December 2008  Communications of the American Ceramic Society  4131  Table I. Mechanical Properties of TaC-Based Ceramics  Material  Relative density (%)  Young’s modulus (GPa)  Vickers’ hardness (GPa)  Fracture toughness (MPa \\x01 m1/2)  Strength (MPa)  Monolithic TaC TaC-10 wt% TaB2 (21001C)  94  98.6  47277 54377  14.170.2 19.470.6  3.570.2 3.470.1  686797 600785  uct.15 Therefore, the oxidation of TaC was much faster than that of TaB2. The mass gain for TaC at 15001C was 85 mg/cm2, which was equal to the theoretical weight gain (14.5%) for the  conversion of TaC to Ta2O5 with all of the carbon lost as CO gas. The TGA results are consistent with the plateau in the mass  gain for TaC in Fig. 4. The oxidation resistance of TaC was in creased slightly by adding TaB2 particles. TaC-10 wt% (11 vol%) TaB2 did not show measurable mass gain until a slightly temperature (B9001C) compared with monolithic TaC higher (B8001C), presumably because of the formation of B2O3 in the TaB2 phase, which acted as a barrier for oxygen transport. Further, the mass gain remained lower for TaC-10 wt% (11 vol%) TaB2 at temperatures below 13501C. Both TaC and TaC-10 wt% (11 vol%) TaB2 oxidized to completion around 15001C.  IV.  Conclusions  TaC ceramics containing 10 wt% (11 vol%) TaB2 were hot pressed to 498.6% relative density at 21001 and 22001C, which  was much higher  than the relative density of monolithic TaC  (85%) hot pressed at  the  same  temperatures. The  enhanced  densification of TaC was due to the physical pinning effect of  TaB2. Using a small amount of a reducing agent such as C or B4C to remove the surface oxide impurities, in addition to TaB2, may enable achievement of full density. The mechanical prop erties of TaC-10 wt% (11 vol%) TaB2 were better than or comparable to those of monolithic TaC due to the higher relative  density of the TaC-10 wt% (11 vol%) TaB2. Young’s modulus increased from 472 GPa for monolithic TaC to 543 GPa for the  TaC-10 wt% (11 vol%) TaB2 composition. Vickers’ hardness also increased from 14.1 to 19.4 GPa. Fracture toughness was similar for both monolithic TaC (3.5 MPa \\x01 m1/2) and TaC-10 (3.4 MPa \\x01 m1/2). The wt% (11 vol%) TaB2 average flexure strength for TaC-10 wt% (11 vol%) TaB2 (600 MPa) was comparable to monolithic TaC (686 MPa). The oxidation resistance  of TaC-10 wt% (11 vol%) TaB2 was marginally improved, compared with TaC, due to the higher oxidation resistance of  the TaB2 phase. Overall, TaB2 proved to be an effective densification aid for TaC based on the relative density, grain size, and mechanical properties of ceramics hot pressed at 21001C.  Acknowledgments  The authors would like to thank Dr. Eric W. Bohannan for the XRD and TGA  analyses, Shi C. Zhang for technical discussion, and Daniel Brenneman for tech nical assistance.  References  1X. Zhang, G. E. Hilmas, W. G. Fahrenholtz, and D. M. Deason, ‘‘Hot Pressing  of TaC With and Without Sintering Additives,’’ J. Am. Ceram. Soc., 90 [2] 393-  401 (2007). 2S. Scholz, ‘‘Some New Aspects of Hot Pressing of Refractories’’; pp. 293-305 in  Special Ceramics 1962, Proceedings of a Symposium held by the British Ceramic  Research Association, Edited by P. Popper. Academic Press Inc., New York, NY,  1963. 3E. Roeder and M. Klerk,  ‘‘Studies With the Electron-Beam Microanalyzer on  Hot-Pressed Tantalum Carbide Having  Small Additions  of Manganese  and  Nickel,’’ Z Metalkd, 54, 462-70 (1963). 4D. C. Halverson, A. J. Pyzik, I. A. Aksay, and W. E. Snowden, ‘‘Processing of  Boron Carbide-Aluminum Composites,’’  J. Am. Ceram. Soc.,  72  [5]  775-80  (1989). 5J. J. Fischer,  ‘‘Hot-Pressing Mixed Carbides of Ta, Hf, and Zr,’’ Ceram. Bull.,  43 [3] 183-5 (1964).  10 µm   Fig. 3. Fracture surface of ride hot pressed at 21001C.  tantalum carbide-10 wt% tantalum dibo vol%) TaB2 was 600 MPa, with a maximum of 765 MPa and a minimum of 467 MPa. The strength values are comparable to  the strength of  the monolithic TaC (686 MPa on average). A  typical fracture surface from the flexure test bars (Fig. 3) reveals  predominantly intragranular with less intergranular fracture for  TaC-10 wt% (11 vol%) TaB2, while monolithic TaC shows exclusively intergranular fracture.13  (4)  TGA Analysis  The oxidation behavior of TaC-10 wt% (11 vol%) TaB2 was evaluated using TGA in air up to 15001C and compared with  TaC and TaB2 ceramics (Fig. 4). The TaB2 ceramic showed the lowest weight gain (20 mg/cm2 at 15001C), probably due to the  formation of a continuous protective layer consisting of Ta2O5 and/or B2O3, which acted as a barrier to the penetration of oxygen at temperatures below 14001C. Similar behavior has been observed in previous studies of oxidation of ZrB2 and HfB2.14 Above B14001C, the evaporation of B2O3 became significant and TaB2 started to oxidize more rapidly. Oxidation of TaC in this temperature regime has been shown to form a nonprotective  Ta2O5 scale due to the evolution of gaseous CO as a by-prod Fig. 4. Mass gain in air for tantalum carbide (TaC), tantalum diboride  (TaB2), and TaC-10 wt% TaB2 by thermo-gravimetric analysis.  \\x0c', '4132  Communications of the American Ceramic Society  Vol. 91, No. 12  6P. Schwarzkopf, R. Kieffer, W. Leszynski, and F. Benesovsky, Refractory Hard  11H. L. Brown, P. E. Armstrong, and C. P. Kempter,  ‘‘Elastic Properties of  Metals: Borides, Carbides, Nitrides, and Silicides, pp. 9. The Macmillan Company,  Some Polycrystalline Transition-Metal Monocarbides,’’ J. Chem. Phys., 45 [2]  New York, NY, 1953. 7S. S. Ordan’yan, V. I. Unrod, V. S. Polishchuk, and N. M. Storonkina,  ‘‘Re 547-54 (1966). 12R. A. Cutler,  ‘‘Engineering Properties of Borides’’; pp. 797-8 in Ceramics and  actions in the System TaC-TaB2 Translated from Poroshkovaya Metallurgiya, 165 [9] 40-3 (1976). 8X.  ‘‘Synthesis, Dens Fahrenholtz,  Zhang, G.  E. Hilmas,  and W. G.  ification, and Mechanical Properties of TaB2,’’ Mater. Lett., 62 [27] 4251-3 (2008). 9P. Chantikul, G. R. Anstis, B. R. Lawn, and D. B. Marshall, ‘‘A Critical  Evaluation of  Indentation Techniques  for Measuring Fracture Toughness:  II,  Strength Method,’’ J. Am. Ceram. Soc., 64 [9] 539-43 (1981). 10A. Bellosi and G. N. Babini,  ‘‘Development and Properties of Ultra-High  Temperature Ceramics—Opportunities and Barriers to Applications’’; pp. 847-64  Glasses: Engineered Materials Handbook, Vol. 4, Edited by S. J. Schneider Jr. ASM  International, Materials Park, OH, 1991. 13X. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Densification and Mechanical  Properties  of TaC-Based Ceramics,’’ Mater.  Sci. Eng. A,  (2008),  doi:10.1016/  j.msea.2008.09.024. 14W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. Zaykoski, ‘‘Refractory  Diborides  of Zirconium and Hafnium,’’  J. Am. Ceram. Soc.,  90  [5]  1347-64  (2007). 15E. L. Courtright, J. T. Prater, G. R. Holcomb, G. R. St.Pierre, and R. A.  in Global Roadmap for Ceramics and Glass Technology, Edited by S. Freiman. John  Rapp,  ‘‘Oxidation of Hafnium Carbide and Hafnium Carbide With Additions of  Wiley & Sons Inc., Hoboken, NJ, 2007.  Tantalum and Praseodymium,’’ Oxid. Met., 6 [5/6] 423-37 (1991).  &  \\x0c']"
},{
  "_id": 38,
  "PDF": "Design, fabrication and high velocity oxy-fuel torch tests of a Cf-ZrB2- fiber nozzle to evaluate its potential in rocket motors.pdf",
  "Text": "['Materials and Design 109 (2016) 709-717  Contents lists available at ScienceDirect  Materials and Design  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t d e s  Design, fabrication and high velocity oxy-fuel torch tests of a Cf-ZrB2fiber nozzle to evaluate its potential in rocket motors  D. Sciti a,⁎, L. Zoli a, L. Silvestroni a, A. Cecere b, G.D. Di Martino b, R. Savino b  a ISTEC-CNR, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, (RA), Italy b Department of Industrial Engineering - Aerospace Division, University of Naples “Federico II”, Piazzale Tecchio 80, I-80125 Naples, Italy  H I G H L I G H T S  G R A P H I C A L  A B S T R A C T  • A segmented nozzle composed of divergent and convergent graphite parts and of a ceramic throat was designed. • The ceramic throat with composition 50 vol% Carbon fiber 50 vol% ZrB2 was machined by EDM from a sintered pellet. • The nozzle was tested in a high velocity oxy-fuel torch (HVOF) in conditions simulating typical exhaust engine flows. • No appreciable erosion of the throat was observed after test at 2.5 Mach supersonic flow and 2730 K flame temperature.  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 8 March 2016 Received in revised form 18 July 2016 Accepted 18 July 2016 Available online 20 July 2016  Keywords:  ZrB2 Carbon fibers Oxidation Ablation Nozzle  1. Introduction  The resistance to ablation of a ceramic nozzle made of a 50 vol% Carbon fiber-50 vol% ZrB2 composite was evaluated in a high velocity oxy-fuel torch (HVOF) in conditions simulating typical exhaust engine flows. The composite was prepared by hot pressing a mixture of ZrB2 powders and C chopped fibers and characterized in terms of microstructural features, strength and toughness. Then, the sintered pellet was machined by electro-discharge machining to obtain a ceramic throat. The throat was assembled to convergent and divergent graphite parts to compose a segmented nozzle and tested in HVOF torch under a heat flux of 2.5 MW/m2 and flame temperature of 2730 K. Fluid dynamic simulations enabled to rebuild the heat field temperatures of the jet flow and of the solid nozzle. The throat well survived 30 s undergoing little oxidation of the frontal zone without dimension or shape variation. No appreciable ablation of the throat was measured. © 2016 Elsevier Ltd. All rights reserved.  There is an increasing demand for advanced materials with temperature capability over 2300 K in highly corrosive environments for aerospace applications. Rocket nozzles of solid or hybrid rocket motors must survive critical thermal, chemical and mechanical environments  ⁎  Corresponding author. E-mail address: diletta.sciti@istec.cnr.it (D. Sciti).  http://dx.doi.org/10.1016/j.matdes.2016.07.090 0264-1275/© 2016 Elsevier Ltd. All rights reserved.  produced by high performance solid propellants, with typical flame temperatures exceeding 2700 K and highly corrosive behavior. The interaction of environmental conditions together with the requirement that dimensional stability of the nozzle throat must be maintained makes the selection of suitable materials extremely challenging. In aeronautical and space applications, C/C, C/SiC and SiC/SiC are the materials currently in use [1,2] for different aero-engine parts and hot structures such as thermal protection systems for re-entry and nozzles for propulsion. Silicon-carbide based composites benefit from the  \\x0c', \"710  D. Sciti et al. / Materials and Design 109 (2016) 709-717  formation of a protective SiO2 surface film, an excellent oxidation barrier at temperatures below 1600 °C in an oxygen-rich atmosphere. However, above this temperature the silica begins to soften dramatically and in low-oxygen atmosphere it develops a substantial vapor pressure. In particular, for propulsion with solid propellants, extensive fractions of Si-based phases are not acceptable. On the other hand, carbon fiber or graphite based materials are plagued by severe erosion in harsh combustion environment. If not controlled, erosion may jeopardize flight safety due to uncontrolled thrust imbalance when, for instance, two boosters are fired at the same time. In recent years, UHTCs (including ZrB2, HfB2, ZrC, HfC and TaC) have been extensively investigated as innovative thermal protection systems [3-5], as well as other applications where oxidation and/or erosion resistance at temperature up to 2300 K are required. More recently, the development of suitable techniques to densify and machine UHTCs at reasonable cost has triggered a considerable amount of testing. TaCbased nozzle inserts were tested in small scale rocket engine [6,7], confirming the outstanding erosion resistance displayed by this class of ceramics. However, these tests also confirmed the lack of capability to resist fast thermal gradients and high pressure without premature failure. In turn, this is due to the low fracture toughness and low thermal shock resistance of these materials. To overcome these limits, reinforced UHTC systems have been studied by several authors following different approaches: 1) enriching the matrix of C/C composites with UHTC phases [8-16], 2) incorporating high aspect ratio secondary phases, mainly SiC whiskers, C or SiC fibers into UHTC matrices [17-20]. Following the latter approach, we developed innovative hot pressed/spark plasma sintered UHTCMCs using randomized carbon or SiC fibers embedded into ZrB2, HfB2 or TaC matrices [19,20]. The advantage of this process is the possibility to exploit conventional powder metallurgical methods (such as mechanical mixing) to introduce fibers in the brittle ceramic matrix, gaining at the same time an effective improvement of key properties, such as flaw tolerance and thermal conductivity. In this perspective, the idea of this work was to fabricate a novel Cf-rich ZrB2 composite, able to resist typical conditions of propulsion, in terms of high temperatures, heat fluxes and corrosive gases. In the literature, the ablation or oxidation behavior of materials for aerospace app l ication is often studied by oxyacety lene tests impinging on cy l indrical pellets constituted o f C/C composites enriched with UHTCs phases [8,13, 21,22]. Tang et al., for instance, showed that the add it ion of ZrB2 to a C/C marked ly reduced the mass erosion rate under a 3 .9 MW/m2 heat flux , compared to a pure C/C [8] . A s im i lar find ing was reported by Pau l et a l . who showed a significant improvement of the erosion resistance for C/C composites enriched with HfB2 or HfC [13]. In this work, a ceramic throat was designed and machined from a 50 vol% Carbon fiber-50 vol% ZrB2 composite. The ceramic throat was assembled into a converging-diverging nozzle typically utilized to expand exhaust gases from the combustion chambers of rockets [6], being the converging and diverging parts made of graphite. The nozzle was tested in a HVOF facility, in a mixture of methane, oxygen and air. Different from the oxyacetylene torch, this facility allows the flame to be accelerated to produce a supersonic flame, with proper control of the pressure and uniform temperature profile. Although the HVOF facility does not fully reproduce the conditions (in terms of gas composition, velocity and shear stresses) present inside the rocket motor , it a llows to expose the test art ic le to temperatures and pressures typ ica l of a harsh combust ion env ironment (5 to 10 atm and 2000-2500 K in the throat) , thus permitting a prequalification of materials for propulsion. Computational fluid dynamic simulations were carried out with a turbulent flow solver to investigate the flow distributions and predict the heat flux and temperatures in the most critical parts of the prototype. In parallel, microstructural characterization was carried out to investigate the erosion and oxidation damage.  2. Experimental  2.1. Nozzle design  The test article has the same shape and size of a graphite nozzle typically used in a small scale rocket engine available at the propulsion laboratory of the University of Naples Federico II. The design of the nozzle studied in the present work is based on the De Laval converging - diverging geometry. The converging element is a truncated cone with a 39° angle, connected to a throat with 9.8 mm inner diameter. The final conical element has a diverging angle of 10° and exit diameter of 16.6 mm. The total length of the nozzle is 35 mm. In order to reduce thermal stresses, this nozzle was segmented in three parts having the outer parts, e.g. the converging and diverging conical elements, made of graphite, while the restricted region around the throat (which is subjected to higher thermal and mechanical stresses) consists of inserts made of different materials, including bulk ceramics [6] or reinforced composites (see Fig. 1). With this design, the ceramic component has a size of 15 mm in length and ~ 5 mm in thickness at the throat level (rather than 35 mm and ~ 10 mm of the whole nozzle), thus longitudinal and radial thermal stress are minimized because of reduced thermal gradients. Furthermore, the compositional design of the ceramic throat was based on the criterion to achieve thermal properties that are as close as possible to those of graphite parts. Fig. 2 shows the sketches of the assembled segmented nozzle and of its parts, with corresponding measures.  2.2. Material production and characterization  Commercial powders were used to prepare the ceramic composites: ZrB2 Grade B (H.C. Starck, Germany), Si3N4 Baysind (Bayer, Germany) as sintering aid, and commercial in-house chopped pitch-derived C fibers. The following composition was produced: 46 vol% ZrB2 + 8 vol% Si3N4 + 46 vol% C fibers. The powder/fiber mixture was ball milled for 24 h in absolute ethanol using silicon carbide media. Subsequently, the slurry was dried in a rotary evaporator. Hot-pressing cycle was conducted in low vacuum (~ 100 Pa) using an induction-heated graphite die with an uniaxial pressure of 30 MPa at 2100 K. The bulk density was measured by the Archimedes' method. The microstructure was analyzed on fractured and polished surfaces by scanning electron microscopy (FE-SEM, Carl Zeiss Sigma NTS Gmbh, Oberkochen, DE) and energy dispersive x-ray spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK). From the same sintered pellet, both sample bars for mechanical testing and the ceramic throat were machined by EDM. The fracture toughness (KIc) was evaluated by fracturing single edge notched beams (SENB) in 3-pt bending configuration. The test bars were 25 mm × 2 mm × 2.5 mm (length by width by  Fig. 1. Segmented nozzle with converging and diverging elements in graphite and throat insert.  \\x0c\", 'D. Sciti et al. / Materials and Design 109 (2016) 709-717  711  Fig. 2. Sketches of: a) Assembled segmented nozzle, b) Converging element, c) Throat insert, d) Diverging element. (Measures in mm).  thickness, respectively) and were notched with a 0.1 mm-thick diamond saw; the notch depth was about half of the bar thickness. The specimens were fractured using a fully-articulated steel three point fixture with a lower span of 20 mm using a electro-mechanic uniaxial testing machine (Instron, 6025). The test bars were loaded with a crosshead speed of 0.5 mm/min. The work of fracture (WoF) was also calculated for each fracture toughness specimen. The total work was determined by measuring the area under the load deflection curve. The WoF was then calculated by dividing the total work by the new surface area produced by fracture (twice the cross sectional area of the un-notched portion of the toughness specimens). The 4-pt bending flexural strength (σ) was measured at room temperature and 1800 K in partially protective Argon atmosphere, according to the European standards EN 843-1 and ENV 820-1 by fracturing five chamfered bars with dimensions of 25 mm × 2.5 mm × 2 mm (length by width by thickness, respectively).  2.3. Nozzle machining  After mechanical and microstructural characterization, the ceramic throat was machined from the sintered pellet by EDM, exploiting the good electrical conductivity of the composite. The graphite converging and diverging parts were machined from a commercial graphite piece (Tokai Carbon G330). The correspondence between the design measures and the real component measures was checked by a high accuracy multi-sensor metrology system, (Smartscope Vantage, gp Inc), with a precision of 0.002 mm. With the same system, the inner throat diameter was measured before and after the HVOF test. The surface roughness, Ra, was measured in the inner wall before and after the test by a stylus-type surface roughness tester (Surftest SJ-301, Mitutoyo).  2.4. HVOF tests: Description of facility and experimental set up for the ablation test  The Sulzer-Metco Diamond Jet (DJ) 27000 was selected as the HVOF torch. Fig. 3a reports a schematic of the system and in particular Fig. 3b is a sectional drawing with the dimensions of the convergent-divergent nozzle of the gun. The DJ gun uses a combination of oxygen, fuel and air to produce a high pressure annular flame, which is characterized by a uniform temperature distribution. In this process, the premixed fuel gas (typically methane, propylene or hydrogen) and oxygen are fed from the annular gap to the air cap, where they react to produce high-temperature combustion gases. The exhaust mixture can reach a temperature of 3000 K, very close to the adiabatic flame temperature of the oxygen-methane combustion. The flame is accelerated by a converging/diverging nozzle to produce a supersonic flame. The combustion chamber is cooled by water and air whilst nitrogen is injected in the central inlet nozzle to prevent it from melting. The segmented ceramic/graphite nozzle was mounted on a mobile support and positioned in front of the HVOF torch, at a distance of 70 mm and exposed for a total period of 30 s, see Fig. 4. As for the gas injected through the DJ gun Table 1 summarizes the overall corresponding mass fractions of the involved reacting species with respect to the total gaseous mass flow rate. The gas flow conditions at the torch exit were calculated assuming that:  • an instantaneous equilibrium is reached at the entrance of the HVOF gun after combustion of methane with oxygen and air (according to Table 1 the overall oxygen-methane ratio is almost equal to 4, i.e. stoichiometric);  \\x0c', '712  D. Sciti et al. / Materials and Design 109 (2016) 709-717  Fig. 3. a) Schematic representation of the HVOF system. b) Converging-diverging nozzle of the HVOF system.  •  adiabatic combustion and local through the nozzle are reached; • all gases behave as ideal.  equilibrium during the passage  carried out by SEM/EDS, focusing on degradation mechanisms of matrix, fiber and interface.  Under these hypotheses, given the mass fraction of the reactants and the geometrical characteristics of the torch nozzle, the HVOF operating conditions in the combustion chamber and at the torch exit can be calculated utilizing the CEA chemical equilibrium code developed by NASA [25]. The results of the described procedure are summarized in Table 2, listing the main combustion products at the torch exit with corresponding molar fraction, and in Table 3, reporting the flow conditions and the gaseous mixture properties. During the experimental test, the surface temperature measurements were carried out using an optical pyrometer focused on the internal surface of the graphite converging element. Post-tests analyses were  Fig. 4. Video capture of the HVOF facility during test.  2.5. Fluid dynamic simulations  Fluid dynamic numerical simulations were carried out to investigate flow conditions and temperature distributions during the ablation test and to evaluate the heat flux on the internal surface of the nozzle, which provides additional information that are difficult to collect experimentally. The simulation was divided in two steps. As mentioned in the previous section, the combustion process in the HVOF was first simulated by means of CEA chemical equilibrium code, which allows to calculate the reacting mixture flow conditions and properties in the combustion chamber and through the torch nozzle. In particular, the assumption of local chemical equilibrium was made for the flow in the combustion chamber and in the torch nozzle. The last hypothesis is justified by the fact that the characteristic convection time is significantly longer than the characteristic chemical time, at the temperatures and pressures of interest. The results of this calculation were given in input to a fluid dynamic solver in order to simulate the flow field from the torch exit through the test sample, under the assumption of mixture of reacting species in chemical non-equilibrium. In particular an axialsymmetric transient simulation was carried out with the commercial fluid dynamic solver ANSYS Fluent™, using the model of compressible, turbulent, reacting flow. The standard k-ε turbulence model, was used to simulate the turbulence, while the Finite-Rate model for the methane-air two-step combustion reaction was used to simulate the  Table 1 Mass fractions of the species involved in the HVOF torch combustion process.  Mass fraction  CH4  0.117  O2  0.478  N2  0.405  \\x0c', 'D. Sciti et al. / Materials and Design 109 (2016) 709-717  713  ð1Þ  Table 2 Species at the HVOF torch exit.  Species  CO CO2 H2 H2O NO N2 OH O2 Others  Mole fraction (%)  6.9 14.2 3.1 38.5 0.2 35.3 1 0.4 0.4  Table 3 HVOF Conditions.  Total pressure (bar) Total temperature (K)  Static pressure (bar) Static temperature (K) Mach number  Total conditions  Static conditions at HVOF gun exit  Gaseous mixture properties at the torch exit  Mean molecular weight (kg/kmol) Specific heat (J/ kg K) Dynamic viscosity [10-5. kg/(m s)] Thermal conductivity (W / m K)  6.9 2986  0.73 2431 2.12  26:51 4126 8:52 0:62  chemical process, taking into account all the main chemical species resulting from the combustion reactions. The pressure, temperature and chemical species concentrations, computed with the CEA software at the torch exit, were imposed along the corresponding inlet surface (see Tables 2 and 3). Furthermore, the computational domain included also the solid regions corresponding to each part of the segmented  nozzle, in which the energy equation is:  \\x00  ρcp  \\x01  s  ∂T ∂t  ¼ λs∇2 T  where ρ is the density, cp is the specific heat capacity and λ is the thermal conductivity, the subscript s indicates that the quantity refers to the solid regions. With the exception of the external wall, which was not exposed directly to the flow and for this reason was considered adiabatic, the thermal coupling condition was set on the interfaces between fluid and solid domains, that is temperature and heat flux continuities:  T f ;int ¼ T s;int  λ f  ∂T ∂n  ¼ λs  ∂T ∂n  s;int  f ;int  ð2Þ  ð3Þ  where n is the normal direction of the interface, the subscript f indicates the fluid state.  3. Results and discussion  3.1. Microstructural composite  features and mechanical properties of the ceramic  The bu lk dens ity of the final pel let is 3 .4 g/cm3 , wh ich corresponds to a fina l re lat ive dens ity o f 85% cons ider ing the theoretical value ca lcu lated with the rule of m ixtures . The assintered microstructure of the compos ite is shown in F ig . 5 . The carbon fibers were homogeneously distributed in the ZrB2 , matrix with a random or ientation , Fig . 5a . Around the carbon fiber , light grey grains were observed, recognized as ZrC, as well as SiC phases deriving from react ion of carbon pockets detached from the fiber with Si3N4, Fig. 5b [23] . These phases were estimated to be around 5 vo l% and 1 vol% , respectively , by image ana lys is performed on  Fig. 5. SEM micrographs showing a) the overall microstructure with a photo of the as-machined throat in the inset, b) a magnification of the fiber/matrix interface, c) the matrix with residual porosity and carbon debris and d) the fracture surface with fiber pull-out denoting a weak fiber-matrix bond with a typical load-displacement curve in the inset recorded during SENB tests.  \\x0c\\x0c\\x0c\\x0c \\x0c\\x0c\\x0c\\x0c \\x0c', '714  D. Sciti et al. / Materials and Design 109 (2016) 709-717  Fig. 6. a) Photo of the nozzle after the HVOF test, b) optical microscopy image of the front part of the nozzle, c) SEM image showing the microstructure of the boxed area in b), d) cross section of a chip showing the oxidized layer, e) microstructure detail of the oxidized layer in d) with the corresponding EDS spectrum inset.  po l ished SEM images . It is apparent that the matr ix contains a significant fraction of porosity, around 15%, despite the application of mechanical pressure at high temperature, Fig. 5c. On the fracture, fiber pu l l-out was observed , wh ich ind icates a weak fiber/matrix interface, Fig. 5d. As for the mechanical properties, the 4-pt. flexural strength was 103 ± 3 MPa at room temperature and increased to 124 ± 25 MPa when measured at 1800 K in Argon atmosphere . The frac ture toughness evaluated through SENB in flexure, was 2.4 MPa·m0.5. In spite of this relatively low value , the load-displacement curves of the notched spec imens disp layed a non-br itt le behav ior , as the example shown in the inset of Fig. 5d. After the first pop-in, a small load drop followed and the curve assumed a tailed shape, typical of stab le crack propagat ion . The work of fracture , WoF , represented by the area underly ing the curves d iv ided by the doub le o f the around 160-180 J/m2 . projected rea l surface [24] , gave va lues These WoF va lues are sim i lar to prev ious ly produced cont inuous fiber-reinforced composites [25] and tw ice those of sim i lar composites containing short carbon fiber, but in lower amount [19]. In the inset of Fig. 5a a photograph of the machined nozzle is reported. The throat final surface roughness, Ra, ranged from 0.6 to 0.64 μm, along the internal profile. The minimum throat width as measured by the Smartscope optical system, was 9.594 ± 0.002 mm.  3.2. Nozzle inspection after HVOF test  The ceramic throat was removed intact from the converging/diverging parts for visual and microscopic inspection. Visual inspection revealed a change in color from grey to white on the frontal part of the ceramic throat, due to surface oxidation, Fig. 6 a,b. This whitish color was also observed inside the throat. On the back side, the color change was less evident, consistently with the lower temperature experienced. As already observed, the most marked effect of the gas flux is a superficial oxidation of the throat walls. ZrB2 oxidizes to ZrO2, carbon fibers placed on the surface are consumed leaving voids, Fig. 6c, SiC  particles leave SiO2-discontinuous residuals, Fig. 6e. On the cross section of a thin oxidized chip, it was possible to observe that the ZrO2 thickness was around 100 μm in the frontal region of the ceramic throat, Fig. 6d,e. Assuming that the thickness of the oxide is proportional to the temperature experienced, it is reasonable to suppose that the thickness inside the throat is even lower than that value. After the test, the measured diameter in the narrowest point was 9.590 ± 0.002 mm, e.g. not significantly changed compared to the initial value (9.594 ± 0.002 mm). The internal roughness increased, due to oxidation, from 0.60-0.64 to 0.75-0.80 μm.  3.3. Fluid dynamic results  For solving Eqs. 1-3, basic properties of the involved materials are needed, e.g. heat capacity and thermal conductivity. In particular, for the converging and the diverging elements the state of the art properties of Tokai Carbon G330 graphite [7] were considered. For the ZrB2Cf composite, we could not directly measure the thermal conductivity and specific heat capacity. Thus, these values were calculated with the rule of mixtures considering the final composition: 0.425 ZrB2 + 0.425 Cf and 15 vol% porosity. The presence of secondary phases, SiC and ZrC, was neglected due to their low amount. For each component, the values available in literature corresponding to the temperature reached  Table 4 Materials properties set in the numerical simulation.  Simulation 1  Simulation 2  Graphite+ Ceramic throat  1800  104  710  3400  90  1100  Graphite+ Ceramic throat covered with 100 μm ZrO2 layer 3400  1800  104  710  2  650  Bulk density ðkg=m3 Þ Thermal conductivity  ðW = m K Þ  Specific heat ð J= kg K Þ  \\x0c', 'D. Sciti et al. / Materials and Design 109 (2016) 709-717  715  Fig. 7. a) Computed Mach number distribution and b) temperature distribution (K) in the flow field at time equal to 15 s.  during the experimental test were considered for the material properties. For ZrB2 thermal conductivity the value of 67 W/mK [26] was used, whilst the thermal conductivity of pitch derived C fibers  considered in this work was 300 W/mK. However, thermal conductivity of carbon fibers can be significantly different in the axial or in the transverse direction [27]. In our case, being the fibers randomly oriented, we  Fig. 8. Maximum temperature profile on the throat insert and the converging surfaces as a function of time.  \\x0c', '716  D. Sciti et al. / Materials and Design 109 (2016) 709-717  nozzle during the test, respectively, calculated with the numerical fluid dynamic model. The typical structure of an over-expanded jet can be observed at the gun exit. The maximum Mach number achieved was around 2.5, immediately after a shock discontinuity which occurs in the diverging element of the nozzle; the highest temperature of the flow field was around 2730 K in the core of the throat. The results of the two numerical simulations (with and without the ZrO2 layer) were very similar: the effect of the zirconia layer is that, due to its low thermal conductivity, the temperature grows faster in the throat inner surface and reaches a maximum value slightly higher, with a difference between the maximum temperatures in the two cases of about 30 K. As a consequence, the surface heat flux along the throat inner surface decreases more rapidly. In the following, only the numerical results in presence of the zirconia layer are presented. Fig. 8 shows the temporal profile of the maximum temperature over the inner surfaces of the throat insert, which increases quickly in the first few seconds and then it tends to reach a steady value . Fig. 9 shows the temperature distribution in the nozzle after 30 s . The temperature over the internal sur face of the throat insert reaches a maximum va lue of about 1540 K after 12 s , that is very close to the experimentally measured value; in the final part of the diverging element, the minimum value of the temperature is instead equal to 1160 K. Finally, Fig. 10 reports the trend of the computed wall heat flux distribution along the inner surface of the ceramic insert (e.g. the most critical part) at different times during the test. It can be noticed that the flux increases in the converging part, reaching a local maximum in correspondence of the throat section. Moreover, at the beginning of the the heat flux up to 3.7 MW/m2 are predicted test, higher values of is cold. The first 5 s are the most critical for the because the material large stresses due to the material thermal expansion. Increasing the material temperature the heat fluxes decrease, but at the end of the test they are still in the order of 1.5 MW/m2.  4. Conclusions  A novel ZrB2 ceramic with 50 vol% randomly oriented C fibers was hot pressed and characterized from the mechanical and microstructural point of view. From the ceramic pelllet, a throat insert was machined in order to fit into a segmented nozzle with converging and diverging graphite elements and exposed to a combination of oxygen, methane and air in a High Velocity Oxygen Fuel facility (HVOF). A fluid dynamic numerical simulation enabled to investigate the flow conditions and the temperature distributions during the experiment. A maximum Mach number of 2.5 and temperature of 2730 K were achieved in the flow field inside the nozzle. The computation of the temperature distribution in the solid part revealed the ceramic  Fig. 9. Computed temperature distribution (K) in the solid region at t = 30s.  hypothesized in first approximation that only half of the fiber volumetric percentage can conduct heat in the radial direction. The calculation yields a value of about 90 W/mK for the composite, which is not very different from the thermal conductivity of graphite (~ 100 W/mK). For the specific heat, a value of 0.65 J g− 1 K− 1 was considered for ZrB2 [26], while for C fiber a typical value of 2 J g−1 K− 1 was taken [28]. Applying the same rule of mixture, a value of 1.1 J g− 1 K− 1 was obtained for the composite. On the basis of the post-test nozzle inspection, two different conditions were simulated. First, we considered the thermal properties of the sintered ZrB2-Cf composite neglecting any surface effect. A second simulation was carried out to account for the oxidation occurring on the throat surface. It was assumed that the ZrB2-Cf throat was uniformly covered by a 100 μm thick surface layer of porous ZrO2 in agreement with the microstructural observations. In this case, we took 2 W/mK as the thermal conductivity of ZrO2, [29] and 650 J g− 1 K− 1 as the heat capacity at 1500 °C [30] (see Table 4, collecting relevant parameters for the two simulations). During the experimental test with the HVOF gun, temperatures as high as 1550 K were measured by the pyrometer on the internal surface of the converging element of the nozzle. Fig. 7a and b show the Mach number and the temperature distribution in the flow field around the  Fig. 10. Total wall heat flux along the internal surface of the nozzle.  \\x0c', 'D. Sciti et al. / Materials and Design 109 (2016) 709-717  717  insert to be the most heavily stressed, with temperature of 1540 K reached on the surface of the throat narrowest part. Although a light oxidation took place in the frontal side of the ceramic insert, it maintained the pristine dimensions and the high volumetric amount of fiber was beneficial in improving the capability of ZrB2 to resist high thermal stresses. No appreciable ablation was observed by measuring the inner throat diameter before and after the test. These first results suggest that this composite, in the segmented nozzle geometry, could be effectively used as a component in hybrid or solid rocket motors and guarantee a high flux stability during combustion.  Acknowledgements  D. Dalle Fabbriche and C. Melandri from ISTEC are gratefully acknowledged for hot pressing cycles and mechanical tests. We wish to thank G. Andalò srl for EDS machining and careful measurements of the tests article dimensions and surface roughness.  References  [2]  [5]  [1] Krenkel W. (Ed.), Ceramic Matrix Composites: Fiber Reinforced Ceramics and Their Application. Wiley-VCH. Verlag:Weinheim; 2008. ISBN: 978-3-527-31361-7. F. Breede, R. Jemmali, H. Voggenreiter, D. Koch, Design and Testing of a C/C-SiC Nozzle Extension Manufactured Via Filament Winding Technique and Liquid Silicon Infiltration, in Design, Development, and Applications of Structural Ceramics, Composites, and Nanomaterials: Ceramic Transactions, John Wiley & Sons, Inc., Hoboken: NJ, USA, 2014 244, http://dx.doi.org/10.1002/9781118889770.ch1. [3] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory Diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364, http://dx.doi.org/ 10.1111/j.1551-2916.2007.01583.x. [4] X. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Densification, mechanical properties, and oxidation resistance of TaC-TaB2 ceramics, J. Am. Ceram. Soc. 91 (2008) 4129-4132, http://dx.doi.org/10.1111/j.1551-2916.2008.02780.x. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, Synthesis, processing and properties of TaCTaB2-C ceramics, J. Eur. Ceram. 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Weng, Densification, microstructure and mechanical properties of ZrB2-SiCw ceramic composites, J. Eur. Ceram. Soc. 29 (2009) 2893-2901, http://dx.doi.org/10.1016/j.jeurceramsoc.2009.03.008. L. Silvestroni, D. Dalle Fabbriche, C. Melandri, D. Sciti, Relationships between carbon fiber type and interfacial domain in ZrB2-based ceramics, J. Eur. Ceram. Soc. 36 (2016) 17-24, http://dx.doi.org/10.1016/j.jeurceramsoc.2015.09.026. [20] D. Sciti, S. Guicciardi, L. Silvestroni, Are short Hi-Nicalon SiC fibers a secondary or a toughening phase for UHTCs? Mater & Des 55 (2014) 821-829, http://dx.doi.org/10. 1016/j.matdes.2013.10.019. [21] X. Zhang, Z. Chen, X. Xiong, R. Liu, Y. Zhu Y. Morphology and microstructure of ZrB2- SiC ceramics after ablation at 3000 °C by oxy-acetylene torch, Ceram, Int 42 (2016) 2798-2805, http://dx.doi.org/10.1016/j.ceramint.2015.11.012. F. Zuo, L. Cheng, L. Xiang, L. Zhang, L. Li, Ablative property of laminated ZrB2-SiC ceramics under oxyacetylene torch, Ceram. Int. 39 (2013) 4627-4631, http://dx. doi.org/10.1016/j.ceramint.2012.10.179. S. Gordon, B.J. McBride, Computer program of complex chemical equilibrium compositions and applications, NASA Reference Publication 1311 (1994). R.L. Smith III, M. JJ Jr., F. SW, Estimation of fracture energy from the work of fracture Int. J. Fract. 156 (2009) 97-102, and fracture surface area: I. Stable crack growth, http://dx.doi.org/10.1007/s10704-009-9350-7. L. Silvestroni, D. Sciti, G.E. Hilmas, W.G. Fahrenholtz, J. Watts, Effect of a weak interface in ZrB2 reinforced with long SiC fibers, Mater. Des. 88 (2015) 610-618, http:// dx.doi.org/10.1016/j.matdes.2015.08.105. J.W. Zimmermann, G.E. Hilmas, W.G. Fahrenholtz, R.B. Dinwiddie, W.D. Porter, H. Wang, Thermophysical properties of ZrB2 and ZrB2-SiC ceramics, J. Am. Ceram. Soc. 91 (2008) 1405-1411, http://dx.doi.org/10.1111/j.1551-2916.2008.02268.x. [27] M. Grujicic, C.L. Zhao, E.C. Dusel, D.R. Morgan, R.S. Miller, D.E. 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},{
  "_id": 39,
  "PDF": "Determination of Retained B2O3 Content in ZrB2‐30 vol_ SiC Oxide Scales.pdf",
  "Text": "['Determination of Retained B2O3 Content in ZrB2-30 vol% SiC Oxide Scales  Kathleen Shugart,*,†,1 Siying Liu, Forrest Craven, and Elizabeth Opila*  Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22904  The  composition of  the borosilicate glass  layer  formed during  oxidation  of ZrB2-30 the extent of B2O3 retention in the oxide during high-temperature oxidation. Oxidation was conducted in stagnant air at 1300°C, 1400°C, 1500°C for  vol% SiC was  determined  to  elucidate  and  times  between  100  and  221 min. Specimens were characterized using mass change and  scanning electron microscopy. After oxidation,  the borosilicate  glass  layer was dissolved from the specimens  sequentially with  deionized H2O and HF acid, using inductively coupled plasma optical emission spectrometry.  to analyze the glass composition  It was  found that  the average B2O3 content to 47 mol%. Retained B2O3 the glass decreased with increasing temperature,  in the glass  scale  ranged  from 23  content  in  the  bulk of  con firming increased volatility with temperature. Boron depth pro files were also obtained in the near  surface region using X-ray  photoelectron spectroscopy and energy dispersive spectroscopy.  The measured B concentrations were  used  to  estimate  the  B2O3 borosilicate  concentration profile and B diffusion coefficients  in the  glass.  Implications  for  the ZrB2-SiC oxidation  process are discussed.  I.  Introduction  O XIDATION  resistance  is  of  great  importance  to Ultra High-Temperature Ceramics (UHTCs), as they are pro posed as Thermal Protection Systems  for hypersonic flight  vehicles. The vehicle leading edges may experience temperatures as high as 2000°C in oxidizing atmosphere.1,2 Studies  on  the  oxidation  characteristics of the UHTC ZrB2 the mid-1960s3-7 with the intent improving the oxidation resistance of this class of materials.  have  been  conducted  since  of  The addition of SiC to ZrB2 rial’s oxidation resistance at  is known to increase the mate higher  temperatures  through  the  formation of a protective borosilicate glass  layer.2,4,8-12  A detailed understanding of  the phases  formed upon oxida tion,  their  composition,  and  the  kinetics  of  formation  are  required as a function of  temperature to enable accurate life  prediction.  The phases present  (in mole  fraction)  after oxidation of  ZrB2-30 vol% SiC are given by the following reaction:  0:61ZrB2 þ 0:39 SiC þ 2:11O2 (g) ¼ 0:61ZrO2 þ 0:39 SiO2 (s/l) þ 0:39 CO(g) þ 0:61B2O3 (l/g)  (1)  where 39 mol% SiC corresponds to the 30 vol% used in this  study. The oxides  form a layered structure, with an outer most  layer of borosilicate glass. A typical example is  shown  in  Fig. 1.  Previous  studies  have  suggested  that ZrB2-SiC  recession rates are controlled by diffusion of oxygen through  the outer glass phase to the base material/oxide interface where the oxidation reactions occur.4,13,14 The B2O3 to SiO2 ratio predicted by Equation 1 is insufficient to define the  glass  composition due  to significant differences  in the vapor  pressures of the two oxides. The vapor pressure of B2O3 at 1500°C is 233 Pa while SiO2 is 3 9 10 suggesting selective volatilization of B2O3 and enrichment of SiO2 in the borosilicate layer relative to the amounts predicted by Equa \\x004 Pa,15,16  tion [1]. Possible B2O3 volatilization reactions include:  B2O3 ð1Þ ¼ B2O3 (g)  (2)  1  2  B2O3 ð1Þ þ 1 2  H2O(g) ¼ HBO2 (g)  (3)  Small changes in B2O3 concentration are shown to have a dramatic effect on oxygen transport in borosilicate glasses. At 1000°C, as little as 1-3 mol % B2O3 in SiO2 on a Si substrate can increase the oxidation rate of Si by two to three orders of magnitude.17 Thus, an understanding of  the B2O3 to predicting  content  in the borosilicate glass  layer  is critical  the oxidation rate of ZrB2-SiC. Oxygen transport in borosilicate glass may occur by per meation or network exchange. Permeation occurs by trans port of molecular oxygen through interstitial  spaces  in the  amorphous borosilicate structure. Network exchange,  similar  to  lattice  diffusion  in  crystalline  phases may  also  occur,  whereby the oxygen jumps from one network site to the next  in the borosilicate glass. Oxygen transport  through pure sil ica has been studied in some detail. Permeation of oxygen in  thermally grown SiO2 on silicon has been shown by Deal and Grove.18 In addition, Ogbuji and Opila19 have found  that  the  enthalpy for  the oxidation of SiC is  essentially the  same  as  that  for Si,  indicating  that  the  controlling mecha nism of oxidation is  the  same  for both materials. Network  diffusion is shown to be much slower than permeation during  oxidation of  silica-forming materials and in silicate glasses, increases.20-22 exchange was shown to  though  the  difference  lessens  as  temperature  A transition  from permeation occur above 1200°C et al.,23,24 while for vitreous silica, Kalen et al. found that both network and interstitial mechanism were in operation.25  to  for  single-crystal  SiC  by  Zheng  However,  few studies have been conducted at  temperatures  as high as those studied here, nor on borosilicate glasses.  The diffusion of B through borosilicate glass  is dependent  on temperature and glass composition. According to Kawaglattice.26  ishi  et al., B diffuses  through Si  sites of  the SiO2 Diffusion rates are found to be on the order of 10 \\x0014 cm2/s, at 10 temperatures ranging between 900°C-1400°C, with low concentrations of B (<15 mol%).26-29 Neither high temperatures nor high B concentrations, have been studied  \\x0020-  for B diffusion in borosilicate glass.  While the specific oxidation rates of ZrB2 and ZrB2-SiC have been measured and semiquantitatively modeled,5-7,14,30-34 a detailed and quantitative understanding of the presence/con centration of B2O3 in the glass layer formed upon oxidation of ZrB2-SiC is currently lacking due to the low sensitivity of most  M. Cinibulk—contributing editor  Manuscript No. 35202. Received June 23, 2014; approved August 27, 2014.  *Member, The American Ceramic Society. 1UES, Inc.; Dayton, OH 45432.  †  Author to whom correspondence should be addressed. e-mail: kns9a@virginia.edu  287  J. Am. Ceram. Soc., 98 [1] 287-295 (2015)  DOI: 10.1111/jace.13236  © 2014 The American Ceramic Society  Journal  \\x0c', 'characterization  techniques  to  light  elements.12,15  Limited  information in the literature supports the presence of B in the  oxide scale and surface depletion. Tripp et al. showed the pres ence of B in the glass phase using Electron Microprobe Analy1300°C and 1400°C.4  sis  (EMPA)  on material  oxidized  at  Karlsdottir et al. discuss the flow of B2O3 to the surface of the oxide, which they analyzed using EPMA and Cathodoluminscence (CL).35,36 They also stated that the surface was depleted  in B2O3 due to its high vapor pressure. This depletion has been shown at the oxide surface by Rezaie et al. using Secondary Ion Mass Spectrometry (SIMS).37 Previous work showed the  formation of borosilicate glass at the very initial stages of oxide  formation using X-ray Photoelectron Spectroscopy though the exact composition was not quantified.38  (XPS),  The  aim of  this  research  is  to  confirm the  presence  of  B2O3 in established oxide scales, in the glass layer formed during oxidation at temperatures of 1300°C-1500°C, and to characterize the compositional gradi to quantify its concentration  ents  formed due to B2O3 volatility. This to the development of a more quantitative description of  information is criti cal  the  oxidation  kinetics  and may  lead  to  approaches  for  improving the currently poor oxidation resistance of ZrB2-30 vol% SiC.  II.  Experimental Procedures  ZrB2-30 vol% SiC bars were provided by Dr. Fahrenholtz (Missouri University of Science and Technology).30,39,40 The  specimens were fabricated using attrition milled powders then hot-pressed. WC contamination (~2 wt%) was observed due to the attrition milling, which used WC  which were  milling media in polyethylene jars. Specimens were cut  from  a  series of billets using an automated surface 40 mm 9 4 mm 9 3 mm and were  grinder  into  bars  of  finished  using  1200 grit diamond abrasive. Bars were sectioned using a low speed diamond saw into coupons approximately 6.2 mm in  length.  Following  sectioning,  each  specimen was  cleaned  ultrasonically  using  detergent  in  deionized  (DI) H2O, DI H2O without detergent, acetone and finally ethanol for 2 min each. Prior to oxidation, the surface area and mass of each  specimen were measured.  Oxidation tests were  conducted in a box furnace  (Rapid Temp, CM Furnaces Inc., Bloomfield, NJ) with molybdenum  disilicide  heating  elements  under  stagnant  air  conditions.  Each coupon was placed on an arc of yttria-stabilized zirco nia  (YSZ, Ortech Inc., Sacramento, CA)  tube  to facilitate  specimen placement and removal  in the furnace, and to limit  contamination from the furnace itself, as previously described.31,38 The arc geometry was chosen to limit contact  between the  specimen and the YSZ tube. Specimen place ment  in the furnace was conducted after the furnace reached  and maintained the desired temperature. Opening the furnace  door dropped the operating temperature approximately 200°C, but the temperature climbed back to within 10°C of  the set point  in 10 s or  less,  leveling out at  the set  tempera ture within 5 min. Specimens were removed from the furnace  while  still  at  temperature,  allowed  to  cool  approximately  15 min after oxidation, and then weighed and measured. The  YSZ boats were also weighed before and after exposure and  negligible mass change was observed.  A baseline  oxidation  test was  conducted  at  1500°C for  100 min. Parabolic oxidation kinetics were assumed based on prior work31,38  in which mass  change  follows  the  relation ship:  Dm  SA  ¼  p ffiffiffiffiffiffiffi  kp t  (4)  where  t  is  the  exposure  time, Dm is  the  oxidation mass  change, SA is  the measured surface area, and kp is sured parabolic oxidation rate constant. The exposure times 1500°C temperatures of  the mea necessary  to  create  the  same mass  gain  as  the  100 min specimen were 1300°C and 1400°C by mg2 cm4/h,  determined  for  using  kp on  values of 5.19 and 4.05 prior work.31 Exposure 1300°C and exposed at these time  respectively,  based  times of 1400°C,  221  and 128 min were  calculated for  respectively. Specimens were  and temperature  conditions  to provide  three specimens with  approximately  the  same mass  gain (and approximately  the  same oxide  thickness), but different oxidation temperatures. Specimens were also exposed for 221 min at 1400°C and 1500°C to provide a series of specimens exposed for the same time at 1300°C, 1400°C, and 1500°C. Three trials of each of these tests were conducted to evaluate variability of the  results. One bar of material was used for  each of  the  three  trials for consistency.  Following mass change measurements, each specimen was  placed in a test tube filled with 15 mL DI H2O maintained at 35°C for 24 h to remove water-soluble B2O3. The specimen was then removed, dried via evaporation at room tem perature and weighed. The DI H2O solution was retained for analysis. The specimen was then placed in another test tube also filled with 15 mL DI H2O maintained at 35°C for 24 h. The second leaching was conducted to confirm that the bulk  of the water-soluble B2O3 was removed during the first H2O soak.41 Again, the specimens were removed and dried and  the DI H2O solution was retained for analysis. After the second leaching bath, each specimen was placed in an HF acid solution (48%-51% HF in H2O) at 35°C for 24 h to dissolve remaining SiO2. In the first trial, the specimens were placed in 3 mL of HF solution  for  24 h  and  then  diluted with  12 mL of DI H2O. mens were soaked in a solution of 12 mL of DI H2O plus 3 mL of HF solution at 35°C for 24 h. The specimens were  In the second and third trials,  the speci removed and dried and the HF solutions were retained for analysis. A very minimal, ~0.05 mL, amount of solution was lost at each removal of the specimen and was unavoidable.  The three solutions  for all  specimens were analyzed for B  and/or Si using  Inductively Coupled Plasma Optical Emis sion  Spectrometry  (ICP-OES, Thermo-Scientific, Waltham,  MA), which has a detection limit on the order of 1 ppb for  most  elements. The purity of  the DI H2O gives mental detection limit of 3 ppb for Si. Precise composition  an experi measurements were made  by  comparison  of  the  emission  spectra from each solution to prepared standards of known  concentrations  and  normalization with  an  internal  yttrium  standard. The standards were prepared using solutions of B  and Si with concentration of 1, 10, and 100 ppm. B and Si  quantities were determined in the H2O solutions, but only Si quantities were determined in the HF solutions, since ZrB2 in HF solutions, as observed by Scanning Electron  dissolves  Microscopy  (SEM,  6700F,  JEOL,  Tokyo,  Japan)  and  described  later. The  total mass  of  all  oxidation  products  Fig. 1. Cross-sectional view of ZrB2-30 1500°C for 100 min in stagnant air.  vol% SiC  oxidized  at  288  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '[ZrO2(s), B2O3(l,g), SiO2(s,l), and CO(g)] was adding the observed mass gained from oxidation and the cal calculated by  culated vaporized mass of both B2O3(g) and CO(g) (assuming full oxidation of the C.) The vaporized mass was  together  calculated using the stoichiometric relation from Equation [1]  and  the  total  amount  of  Si measured  from ICP-OES,  as  described in the  results  section. Finally,  each specimen was  carbon coated with the Precision Etching and Coating Sys tem (PECS),  then examined in the SEM and characterized  with  Energy  Dispersive  Spectroscopy  (EDS,  Princeton  Gamma-Tech Inc., Princeton, NJ)  to ensure the borosilicate  glass layer was completely removed.  Chemical analysis by XPS (PhiVersa Probe XPS spectrom eter, Chanhassan, MN) was performed on two additional specimens oxidized at 1500°C for 10 and 100 min to deter mine  surface B and  Si  concentrations  in  the  glass  layer.  Depth  profiling was  conducted  by Ar  sputtering  at  3 kV  energy on a nonspinning specimen. The  compositional data  were collected on the surface and at depths of approximately  50, 100, and 150 nm into the scale. XPS was performed using  monochromated Al Ka radiation of 1486.6 eV, with a bandpass energy of 11.75 eV and an increment of 0.1 eV/step.  The resolution on Ag3d5 is 0.6 eV. The analyzed area each specimen was about 0.5 mm2 and the  for  scans were per formed  under  identical  conditions. Curve  fitting was  done  using the Gaussian-Lorentzian peak and Shirley background  algorithm of the Phi data massaging program MultiPak. EDS (Oxford Instruments Aztec X-MaxN 150, Concord,  MA) was performed on the surface of a specimen oxidized at 1500°C for 100 min using a series of electron beam accelerat ing voltages  (5, 10, and 20 kV) at a constant working dis tance  of  15 mm.  The  change  in  B  composition  with  accelerating voltage and simulated sampling depth was used  to provide additional  information about  the B concentration  profile  in the borosilicate  glass,  as described further  in the  results section.  III.  Results  (1)  B2O3 Concentration in Borosilicate Glass  SEM and EDS results for two of  the specimens after borosil icate  glass  digestion  are  shown  in Fig. 2.  SEM was  per formed  using  a  beam voltage  of  5 kV  unless  otherwise  specified to increase  sensitivity  to elements of  light  atomic  mass during EDS. Comparing the SEM and EDS results of  glass dissolution trials,  it  can been seen that  trial  1  (3 mL  HF, 24 h)  removed all of  the glass and ZrB2 leaving only SiC. Trial 2 used more dilute HF acid to  from the  sur face,  dissolve the glass  layer,  leaving some ZrO2 grains and miniresidual glass on the specimen surface, which were not  mal  found in trial 1. These  results  indicate  that  the assumption  that all  the glass was  removed in the water and HF dissolu tion procedure  is  reasonable,  ensuring the  ICP-OES results  provide an accurate average  composition of  the borosilicate  glass  formed during oxidation. Note that SiC is not  soluble  in H2O or HF as confirmed by SEM (Fig. 2) so that Si content in the solutions must be attributed to SiO2. The specific mass change from the three oxidation trials  and the B2O3 “Mass Change-Oxidation” is the measured change in mass of  leaching process are  reported in Table I. The  the specimen after oxidation under the stated conditions. The  “Mass Change-Leaching”  is  the mass of material  removed  after both of  the H2O soaks. This mass was determined by subtracting the specimen mass after the H2O baths from the expected to be composed of  mass  after oxidation and was  mostly B2O3 with trace amounts of SiO2. The B2O3 retained in the glass layer after oxidation was calculated assuming:  (1)  an  equal  rate  of  consumption  of  the ZrB2 and SiC in the base material during oxidation; the water leaching steps removed all of the B from the layer;41 and (3) all of  (2)  glass  the Si  from the glass  layer was  removed  in  the H2O and HF acid  soaks. The  validity  of  these assumptions is discussed later. For this calculation,  the  concentrations of B and Si  in the glass  layers of each speci men were determined by ICP in mg/L and then multiplied by  0.015 L,  the volume of each solution. This yielded the mass  in mg of dissolved B and Si, which was then easily converted  to moles. Since the SiO2 does not volatilize in this ture range, multiplication of the amount of Si in moles by  tempera the  stoichiometric  ratio  of B to  Si  expected  from Equa tion (1) yields the total amount of B expected if none volatil izes  (see  Appendix  A  for  an  example  calculation.)  Comparison of  the B concentration found from ICP-OES to  the  total  amount oxidized according  to the  above  calcula tions  gives  the  percent  of  retained B. Table II  provides  a  summary of these calculations. Percent B2O3 given by 100%-% retained.  volatilized  is  Figure 3  is  a plot of  the  total mass of oxides  generated  versus  temperature  for  specimens which were oxidized with  the  goal  of  attaining  the  same  amount  of  oxidation. The  total mass of generated oxides includes the mass gained dur ing oxidation and the calculated mass of B2O3(g) and CO(g) which have formed and vaporized. The similar mass results  for the three temperatures and three trials show that  the cho sen times do provide similar quantities of oxidation, allowing  comparison between these  tests. Figure 4  is  a plot of  total  mass of  generated oxides  versus  temperature  for  specimens  all oxidized for  221 min,  showing with 80% confidence  an  increase in the amount of oxide formed with increasing tem perature.  XPS  results  of  oxidized  specimen  surfaces  are  given  in  Fig. 5. For both 10 and 100 min exposures,  less  than 1% B  was evident on the surface, but after sputtering 50 nm,  the B  content  became measureable  and  steadily  increased  with  depth to 150 nm into the  surface,  the deepest  sputter depth  (a)  (b)  Fig. 2.  SEM/EDS results for ZrB2-30 vol% SiC oxidized at 1500°C for 100 min in stagnant air after two DI H2O soaks and dissolution 35°C for in (a) 3 mL of HF solution at 24 h (b) 3 mL of HF solution diluted with 12 mL of DI H2O at 35°C for 24 h, the area with the most products left behind after dissolution.  showing  January 2015  Determination of Retained B2O3  289  \\x0c', 'analyzed. The B content measured by XPS at depths greater  than  50 nm after  100 min  oxidation was  about  5% lower  than  it was  after  10 min  oxidation, while  the  Si  content  increased 5%. The O and Zr values were equivalent and con stant for the two scales.  A boron concentration profile was  estimated using EDS 1500°C for  results  obtained  from a  specimen  oxidized  at  100 min,  at  a  series  of  accelerating  voltages  (5,  10,  and  Table I.  Mass Change Data for ZrB2-30 vol% SiC Oxidized Under Stagnant Air Conditions After Oxidation and After H2O Soak (Leaching)  Condition  Trial #  Mass change - Oxidation (+mg/cm2)  Average (+mg/cm2)  Mass change - Leaching (-mg/cm2)  Average (-mg/cm2)  1500°C 100 min  1  2.81  2.90 \\x06 0.26  0.70  0.91 \\x06 0.28  2  3.12  1.23  3  2.70  0.80  1400°C 128 min  1  3.85  4.55 \\x06 0.68  1.13  2.07 \\x06 1.05  2  5.21  3.20  3  4.60  1.88  1300°C 221 min  1  3.46  4.61 \\x06 1.06  2.39  2.38 \\x06 0.37  2  5.52  2.75  3  4.86  2.00  1400°C 221 min  1  6.28  7.10 \\x06 1.04  4.29  3.94 \\x06 0.45  2  8.27  4.11  3  6.76  3.43  1500°C 221 min  1  4.38  4.91 \\x06 0.49  1.48  2.02 \\x06 0.74  2  5.00  1.71  3  5.36  2.87  Table II.  Summary of B Retained in the Borosilicate Scale After Oxidation of ZrB2-30 vol% SiC Under Stagnant Air Conditions. A High Percentage of the B does not Volatilize under the Conditions Tested, Leading to a High B2O3 Concentration in the Borosilicate Glass  Conditions  Trial #  B2O3 content  of glass (mol%)  Average B2O3  content  (mol%)  % of Total B retained  in glass (mol%)  Average B  retained (mol%)  “Constant”  Oxide thickness  1500°C 100 min  1  27.3  27.3 \\x06 4.1  23.9  22.6 \\x06 4.7  2  29.2  26.4  3  21.3  17.3  1400°C 128 min  1  31.9  36.8 \\x06 5.7  28.5  37.3 \\x06 10.1  2  43.1  48.3  3  35.5  35.1  “Constant”  time  1300°C 221 min  1  43.2  42.9 \\x06 1.6  43.3  46.2 \\x06 4.0  2  44.3  50.8  3  41.1  44.6  1400°C 221 min  1  34.2  36.9 \\x06 2.4  37.3  38.8 \\x06 1.9  2  39.0  40.9  3  37.4  38.1  1500°C 221 min  1  27.3  28.2 \\x06 2.0  24  25.1 \\x06 2.5  2  26.8  23.4  3  30.4  28  Fig. 3.  Total mass of generated oxides versus temperature for ZrB230 vol% SiC specimens oxidized to generate the same oxide thickness. 1300°C specimens oxidized for 221 min, 1400°C specimens oxidized for 128 min, and 1500°C specimens oxidized for 100 min.  Fig. 4.  Total mass of generated oxides versus temperature for ZrB230 vol% SiC specimens oxidized for 221 min.  290  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '20 kV.) The Casino simulation program was used to deter mine the depth of voltage.42 These  the sampling volume for each accelerating  values  are  given in Table III. A series of  layers were  simulated with  the Casino  program using  the  compositional data determined both from XPS and EDS and the density data from Bruckner et al.41 The sampling volume  for  the EDS was determined based on the  layered borosili cate structure. The first  layer of  the final structure was set  to  be  50-nm thick  with  a  composition  of  6.1 mol% B2O3 from averaging the XPS results for the sur (remainder SiO2) face and 50 nm into the borosilicate. This composition has a density of 2.21 g/cm3. The next 100 nm was assumed to have  a composition of 8.8% B2O3 (remainder SiO2) ing XPS results acquired at 50 and 150 nm depth, with a 2.18 g/cm3. The next  from averag corresponding density of  three  layers  were  assumed  to  have  a  composition  of  13.0,  18.7,  and  31.0 mol% B2O3 (remainder SiO2) positions were determined from mol% B and Si measured by  respectively. These  com EDS using  acceleration voltages of  5,  10,  and 20 kV. The  last  three  layer  thicknesses were adjusted using an iterative  process  to determine  the  average  sampling depth  for  each  accelerating voltage. Table IV provides the layer information  used to generate  the Casino simulations. The  three Casino  simulations  of  the  sampling  volume  at  5,  10,  and  20 kV  based on final  iterations of  the layered structures were used  to provide B concentration and depth values. An example  simulation for  5 kV is given in Fig. 6. The final depth for  each  acceleration  voltage was  determined  using  plots  of  depth versus normalized hits  (electrons hitting atoms  in the  material), as given in Fig. 7 for 5 kV. Estimated B concen trations  from the EDS results  are plotted along with XPS  data in Fig. 8.  (2)  Calculated B Concentration Gradient  The borosilicate glass  layer  formed after oxidation of ZrB230 vol% SiC was approximated as a binary, semi-infinite sys tem, and the error function solution to Fick’s 2nd Law was used to calculate the B concentration gradient:43  CB ¼ C1 erfðgÞ  (5)  g ¼  x  4Dt  (6)  erfðgÞ ¼ 2ffiffiffi p  p  Z  g  0  e  \\x00k2  dk  (7)  where CB borosilicate  is  the  concentration of B at distance x into the  glass  from the  gas  interface, C∞ is  the  equilib rium concentration of B at  the oxide/base material  interface,  D is  the diffusion coefficient of B in the borosilicate  glass,  and t is time. The following boundary conditions were used: BC1: CB = 0 at x = 0 BC2: C∞ = 0.289mole fraction (fixed by ZrB2-30 vol% SiC) at x = 20.2 lm These boundary conditions are reasonable as XPS shows a very low B concentration (<1%) on the  stoichiometry of  surface, and O2(g) diffuses through the borosilicate and forms new oxide rapidly  in comparison to B diffusion outward. The concentration of  B at  the borosilicate glass/base material  interface is assumed  constant. Choosing  a fixed time,  the B diffusion coefficient  was determined using:  Fig. 5.  XPS results for ZrB2-30 vol% SiC oxidized at 1500°C for 10 (open symbols) and 100 (closed symbols) min in stagnant air.  Table III.  Concentration Results from EDS Performed at  Varying Accelerating Voltages, Showing Increased B  Concentration Deeper in the Borosilicate Scale  Accelerating  voltage (kV)  B concentration  (at%)  B2O3 concentration  (mol%)  Simulated average  sampling depth (nm)  5  7.6  13.0  220  10  11.0  18.7  770  20  18.5  31.0  2900  January 2015  Determination of Retained B2O3  291  \\x0c', 'DB ¼ ½ð1 \\x00 Cavg t  C1  Þ  ffiffi  p  p  xl  2  \\x8a2  (8)  Solutions were obtained for a specimen oxidized at 1500°C after 100 min, which forms a borosilicate layer 20.2 lm thick  (BC2),  since  ICP-OES, XPS, and EDS compositional  infor mation are available  for  this oxidation condition. The aver age  B  concentration  in  the  borosilicate  glass  after  this  oxidation treatment given by ICP-OES was 15.4 mol%, and  the concentration at  the oxide/base material  interface is assu med to be  equal  to that  given by  equilibrium, 28.9 mol% found to be ~1.2 9 represents an average  (BC2). The B diffusion coefficient was \\x0010 cm2/s. This diffusion coefficient 10 B diffusion coefficient in borosilicate glass resulting from ZrB2-30 vol% SiC oxidation at 1500°C for 100 min. B diffusion coefficients were also calculated which provided best fits \\x0013 cm2/s to the XPS data and the EDS data and were 1 9 10 \\x0012 cm2/s, respectively. The DB values will be disand 5 9 10 cussed below. Resulting B concentration profiles are shown  in Fig. 8.  IV.  Discussion  (1)  B2O3 Concentration in Borosilicate Glass  Assuming  a  linear  concentration gradient  from 0.289 mole  fraction B (0.61 mole  fraction B2O3) at diboride interface (which corresponds to ZrB2-30 vol% SiC) interface, the average  the glass/zirconium  to  zero  at  the  gas/borosilicate  glass  B2O3 mole fraction is 0.305 which is results obtained at 1400°C and 1500°C,  intermediate  to  the  shown in Table II.  This  suggests  that  the  ICP-OES results are  reasonable. The  ICP-OES results 1500°C relative  confirm a 1300°C,  higher  depletion  of  B2O3 expected B2O3 gain behavior of the speci at  to  consistent with  volatility. Comparing  the mass  mens (Table I) to the total mass of oxides generated in Fig. 3  corroborates this B2O3 volatility.  (2)  Calculated B Concentration Gradient  The B diffusion coefficient  is expected to vary with borosili cate  composition  and will  therefore  vary with  position  in  the oxide  scale. The  coefficient determined for  the  average 1500°C  composition  from  ICP-OES after \\x0010 cm2/s. 1.2 9 10 the surface  oxidation  at  for  100 min  was  The  estimated  EDS  data, which  is  nearer  to  and  thus  lower  in B  content, was  best  fit with  a  boron  diffusion  coefficient  of  Table IV. Layer Information Used in Casino Simulations for EDS Penetration Depth Determinations of Borosilicate Glass Formation on ZrB2-30 vol% SiC After Oxidation at 1500°C for 100 min in Stagnant Air in Standard Box Furnace  Layer  Thickness  (nm)  Total depth  (nm)  Measured B  concentration (at%)  Composition  (mol% B2O3)  Assumed density (g/cm3)  Source of B data  1  50  50  0.1, 3.8  6.1  2.21  XPS at surface and 50 nm  2  100  150  3.8, 5.4  8.8  2.18  XPS at 50 nm and 150 nm  3  70  220  7.6  13.0  2.15  EDS, 5 kV  4  550  770  11  18.7  2.12  EDS, 10 kV  5  2130  2900  18.5  31.0  2.06  EDS, 20 kV  Fig. 6.  Casino  simulation  showing EDS  sampling  volume  for  an  accelerating voltage of 5 kV in borosilicate glass formed on ZrB2-30 vol% SiC after oxidation at 1500°C for 100 min in stagnant air in  standard box furnace.  Fig. 7.  Casino simulation of depth versus normalized hits generated  for 5 kV in the simulated borosilicate glass.  Fig. 8. XPS and EDS results 5 lm of 1500°C  showing B concentration in the first  the  borosilicate  glass  scale  for  a  specimen  oxidized  at  for  100 min  in  stagnant  air. Dashed  line  indicates  B  concentration  profile  estimated  using B diffusion  coefficient  from  best fit  to XPS results. Dotted line indicates B concentration profile  estimated using B diffusion coefficient  from best fit  to EDS results.  Solid line indicates B concentration profile estimated using calculated  B diffusion coefficient from ICP-OES results.  292  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '\\x0012 cm2/s. The compositional data from the XPS rep5 9 10 the lowest B concentration due to proximity to the  resents  surface and was best fit with a boron diffusion coefficient of \\x0013 cm2/s. All of 1 9 10 than previous measured values for B diffusion through boro\\x0014 cm2/s when extraposilicate, which are on the order of 10 lated to 1500°C.27,29,44 The high values of the B diffusion  these diffusion coefficients are higher  coefficient estimated in this work are reasonable since boro1500°C as silicate glass with 27 mol% B2O3 is liquid at shown in the SiO2-B2O3 phase diagram (Fig. 9).45 The \\x0013 cm2/s, mated B diffusion coefficient of 1 9 10 obtained from best fit to the XPS data is reasonable as these data were  esti taken near the surface where less B is present. The glass here  will be more viscous and lower diffusion rates are predicted.  Figure 8 compares B concentration profiles based on best fit  for  the  ICP-OES average  composition, XPS data, and EDS  estimates.  Note also that the average B concentration determined by ICP-OES is a lower bound for B concentration. Bruckner41  states that 95% of B2O3 is removed from borosilicate glasses by a warm water leach, however, Adams and Evans46 state  that only 30% is  removed.  It  is not possible to confirm the  extent  of B2O3 ZrB2 dissolves in HF. If B leaching in H2O is incomplete, average B composition in the borosilicate scale would  leaching  in warm water  at  this  time  since  the  be  higher  and  the  calculated B concentration would  increase,  bringing it  in closer agreement with the XPS and EDS gradi ent results in Fig. 8.  A B2O3 concentration gradient proposed in Fig. 10, for the first  in the borosilicate layer  is  time  providing  composi tional  information about  this  layer. The B2O3/SiO2 is assumed to be equal  ratio at  the oxide/substrate  interface  to the  B/Si  ratio  in  the  base material. This  is  justified  by  two  observations. First, XPS analysis  showed that B-O and Si-O  bonds are both observed on the oxidized surface after only  10 s, indicating both oxides are forming simultaneously at the start of oxidation.38 Second, a nearly uniform oxidation  front  is  observed  between  ZrB2/SiC and the oxides, as shown in Fig. 1. The presence of the ZrO2+C layer between  the glass  layer and the base material  is neglected in Fig. 10  and is  assumed to have  little  effect on the B/Si boundary  condition at  the  interior  glass  interface  (K. N. Shugart &  E.  J. Opila,  in preparation). The B2O3 gradient calculated by converting the B concentration gradient  in the plot  was  estimated from ICP-OES in Fig. 8 to oxide. The discontinuborosilicate/ZrB2+SiC interface the average B concentration determined by ICP-OES  ity  at  the  results  from the  use of  to calculate an average diffusion coefficient.  (3)  Implications for Oxidation  It  is confirmed here that considerable B2O3 is retained in the glass layer which has important effects on oxidation, assum ing the rate-limiting step is oxygen diffusion in the borosili1500°C, layer.4,17 At  cate  glass  the  liquidus  curve  lies  at  approximately 95 mol% B2O3, as surface of the glass layer depleted in B2O3 will be composed of a more viscous SiO2 then the interior of the glass layer with higher B2O3 content. Temperature transport will be  seen in Fig. 9. Thus,  the  effects  on  oxygen  very  complex.  Increasing  temperature  depletes the borosilicate glass in B2O3, slowing oxygen transport by compositional effects at the surface, but at the same  time  increases oxygen transport  since  it  is an activated pro cess.  The  apparent  activation  energy  for  oxidation  then  reflects  some combination of B2O3 vaporization, B transport outward due to the driving force induced by the B2O3 concentration gradient, and oxygen transport inward in the non uniform glass structure.  V.  Conclusions  Significant B2O3 scale formed during oxidation of ZrB2-30 vol% SiC at temperatures between 1300°C and 1500°C. Borosilicate composi is  found (27-43 mol%)  in the borosilicate  tions measured using ICP-OES (average value), EDS & XPS  were used to estimate B diffusion coefficients \\x0013 cm2/s to 10 for B2O3-rich and respectively. The variation in  in the borosili cate  ranging  from 10  \\x0010  SiO2-rich portions of B concentration across  the scale,  the oxide  scale  implies variations  in  Tm, DB, and DO through the resulting complex temperature dependence for oxygen perme thickness of  the  scale  and a  ation,  the likely rate controlling mechanism for oxidation of  ZrB2-SiC.  Fig. 9.  B2O3-SiO2 Equilibria Diagrams,  phase  diagram,  from  ACerS  NIST  Phase  figure  11777.  (Reprinted with  permission  of  The American Ceramic Society.)  Fig. 10.  Approximate  B2O3 formed during oxidation of ZrB2-30 vol% specimen oxidized at 1500°C for 100 min  concentration  (mol%)  profile  in  borosilicate  glass  scale  SiC. Values given are for  in  stagnant  air,  using  average  B  concentration  (ICP-OES)  to  determine  the  average  diffusion  coefficient.  A majority  of  the  borosilicate glass is shown to have sufficient B2O3 to be liquid at oxidation temperature.  the  January 2015  Determination of Retained B2O3  293  \\x0c', 'Appendix A  Oxidation reaction: 0.61 ZrB2+0.39 SiC+2.11 O2(g)=0.61 ZrO2+0.39 SiO2(s 0.39 CO(g)+0.61 B2O3(l/g) Example calculation for: 1500° C, 100 minutes, Molar mass for B2O3= 69.62 g/mole Molar mass for CO= 28.00 g/mole  /l)+  trial 1  Si  to B ratio in starting material:  B  Si  ¼ 0:61mol \\x03 2 0:39mol  ¼ 3:13  [A1]  Total Si  in SiO2 measured via ICP-OES:  2:97mg/L þ 0:95mg/L þ 86:05mh/L ffi 89:97mg/L  [A2]  (H2O soak 1 + H2O soak 2 +HF soak= total Si) Measured Si (mol):  89:97mg/L \\x03  0:015L  28:09g/mol  \\x03  1g  1000mg  ¼ 4:80 \\x03 10 \\x005mol  [A3]  Total B (mol)  consumed  during  oxidation,  predicted  from  [A1] and [A3]:  4:80 \\x03 10 \\x005mol \\x03 3:13 ¼ 1:50 \\x03 10 \\x004mol  [A4]  Total B measured via ICP-OES:  22:03mg/L þ 3:91mg/L ¼ 25:94mg/L  [A5]  (H2O soak 1 + H2O soak 2 = total B) Measured B (mol):  25:94mg/L \\x03  0:015L  10:81g/mol  \\x03  1g  1000mg  ¼ 3:60 \\x03 10 \\x005mol  [A6]  Percentage of B retained ([A6]/[A4]):  3:60 \\x03 10 \\x005mol 1:50 \\x03 10 \\x004mol  ¼ 0:24 ¼ 24%  [A7]  Vaporized B (mol):  1:50 \\x03 10 \\x004mol \\x00 3:60 \\x03 10 \\x005mol ¼ 1:14 \\x03 10 \\x004mol  [A8]  Total B2O3 and CO vaporized:  1:14 \\x03 10 \\x004mol  2 \\x03 1000mg 1g  \\x03 69:62g mol  þ 4:80 \\x03 10 \\x005mol \\x03 28g mol  \\x12  \\x13  ¼ 5:31mg  [A9]  Surface area of specimen: ð6:17mm \\x03 3:99mmÞ \\x03 2 þ ð6:17mm \\x03 2:97mmÞ \\x03 2 þ ð2:97mm \\x03 3:99mmÞ \\x03 2 ¼ 109:59mm2  [A10]  Measured mass after oxidation minus measured mass before  oxidation:  384:95mg \\x00 381:87mg ¼ 3:08mg  [A11]  Total mass of generated oxides  (ZrO2+SiO2+B2O3+CO) per  area:  ð3:08mg þ 5:31mgÞ 109:59mm2  \\x03 100  mm2  cm2  ¼ 7:67mg/cm2  [A12]  Acknowledgments  The authors acknowledge  Joe Hagan and Bohuslava McFarland (University  of Virginia)  for  the ICP-OES results and Dr. Wayne Jennings  (Case-Western  Reserve University)  for  the XPS results. The authors would like to thank Dr.  William Johnson for helpful advice and discussions on calculating the boron  diffusion coefficient. The authors would also like  to acknowledge Eric Neu man  and Dr. William Fahrenholtz  at Missouri University  of  Science  and  Technology for  the ZrB2-30 vol% SiC material. by The National Hypersonic Science Center-Materials  Initial  funding for  this work  was  provided  and  Structures.  References  1A. Paul, D. D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. G. P. Bin ner, , B. Vaidhyanathan, A. Heaton, P. Brown, and W. E. Lee, “UHTC Composites for Hypersonic Applications,” J. Am. Ceram. Soc., 91, 22-28 (2012). 2A. Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby, “Oxidation  of ZrB2-SiC Ceramics Under Atmospheric and Reentry Conditions,” Refract. Appl. Trans., 1, 1-7 (2005). 3H. C. Graham, H. H. Davis, I. A. Kvernes, and W. C. Tripp, “Microstruc tural Features of Oxide Scales Formed on Zirconium Diboride Materials,” Mater. Sci. 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Fahrenholtz, “Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,” J. Am. Ceram. Soc., 90, 143-8 (2007). 16S. N. Karlsdottir and J. W. Halloran, “Formation of Oxide Scales on Zir J. Am.  conium Diboride-Silicon Carbide Composites During Oxidation: Relation of Subscale Recession to Liquid Oxide Flow,” J. Am. Ceram. Soc., 91, 3652-8  (2008). 17J. Sclichting, “Oxygen-Transport Through Glass Layers Formed by a Gel Process,” J. Non Cryst. Solids, 63, 173-81 (1984). 18B. E. Deal and A. S. Grove, “General Relationship for dation of Silicon,” J. Appl. Phys., 36, 3770-8 (1965). 19L. U. J. T. Ogbuji and E. J. Opila, “Comparison of the Oxidation Kinetics of SiC and Si3N4,” J. Electrochem. Soc., 142, 925-30 (1995). 20M. A. Lamkin, F. L. Riley, and R. J. Fordham, “Oxygen Mobility in Silicon Dioxide and Silicate Glasses: A Review,” J. Eur. Ceram. Soc., 10, 347-67  the Thermal Oxi (1992).  294  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', 'January 2015  Determination of Retained B2O3  295  21J. D. Cawley, J. W. Halloran, and A. R. Cooper, “Oxygen-18 Tracer the Passive Thermal Oxidation of Silicon,” Oxidation Metals, 28, 1-  Study of  16 (1987). 22J. A. Costello and R. E. Tressler, “Isotope Labeling Studies of the Oxidation of Silicon at 1000°C and 1300°C,” J. Electrochem. Soc. Solid State Sci, 131, 1944-1947 (1984). 23Z. Zheng, R. E. Tressler, and K. E. Spear, “Oxidation of Single-Crystal Silicon Carbide: I, Experimental Studies,” J. Electrochem. Soc., 137, 854-8 (1990). 24Z. Zheng, R. E. Tressler, and K. E. Spear, “Oxidation of Single-Crystal Silicon Carbide. Part II. Kinetic Model,” J. Electrochem. Soc., 137, 2812-6 (1990). 25J. D. Kalen, R. S. Boyce, and J. D. Cawley, “Oxygen Tracer Diffusion in Vitreous Silica,” J. Am. Ceram. Soc., 74, 203-9 (1991). 26K. Kawagishi, M. Susa, T. Maruyama, and K. Nagata, “Boron Diffusion in Amorphous Silica Films,” J. Electrochem. 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Am.  in  Air, Oxygen, and Nitrogen,” California Inst. of Tech., Pasadena Jet Propul sion Lab, No.JPL-PR-20-252, (1955). 34T. A. Parthasarathy, R. A. Rapp, M. Opeka, and R. J. Kerans, “A Model for the Oxidation of ZrB2, HfB2 and TiB2,” Acta Mater., 55, 5999-6010 (2007).  35S. N. Karlsdottir  and  J. W. Halloran,  “Formation  of Oxide Films  on  ZrB2-15 vol% SiC Composites During Oxidation: Evolution With Time and Temperature,” J. Am. Ceram. Soc., 92, 1328-32 (2009). 36S. N. Karlsdottir, J. W. Halloran, and C. E. Henderson, “Convection Pat terns in Liquid Oxide Films on ZrB2-SiC Composites Oxidized at a High Temperature,” J. Am. Ceram. Soc., 90, 2863-7 (2007). 37A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Oxidation of Zirconium Diboride-Silicon Carbide at 1500°C at a Low Partial Pressure of Oxygen,” J. Am. Ceram. Soc., 89, 3240-5 (2006). 38K. Shugart, W. Jennings, and E. Opila, “Initial Stages of ZrB2-30 vol% SiC Oxidation at 1500° C,” J. Am. Ceram. Soc., 97, 1645-51 (2014). 39A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  “High-Strength Zirconium Diboride-Based Ceramics,” J. Am. Ceram. Soc., 87, 1170-2 (2004). 40A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Evolution of Structure  During the Oxidation of Zirconium Diboride-Silicon Carbide 1500 °C,” J. Eur. Ceram. Soc., 27, 2495-501 (2007). 41V. R. Bruckner and J. F. Navarro, “Physikalisch-Chemische Untersuchungen im System B2O3-SiO2,” Glastechn. Ber, 39, 283-93 (1966). 42D. Drouin, A. R. Couture, D. Joly, X. Tastet, V. Aimez, R. Gauvin, 42—A Fast  and Easy-to-use Modeling Tool  “CASINO V2.  in Air up to  for  Scanning 92-101  Electron Microscopy  and Microanalysis  Users,”  Scanning,  29,  (2007). 43J. Crank, The Mathematics  of Diffusion;  pp.  1-414. Oxford University  Press, New York, 1975. 44G. D. Bagratishvili, R. B. Dzhanelidze, D. A. Jishiashvili, L. V. Piskanov skii, and Z. N. Shiolashvili, “Boron Diffusion From a Reactively Sputtered Glass Source in Si and SiO2,” Physica Status Solidi (a), 56, 27-35 (1979). 45T. J. Rockett and W. R. Foster, “Phase Relations in System Boron OxideSilica,” J. Am. Ceram. Soc., 48, 75-80 (1965). 46P. B. Adams and D. L. Evans, “Chemical Durability of Borate Glasses,” Mat. Sci. Research, 12, 525-37 (1978).  h  \\x0c']"
},{
  "_id": 40,
  "PDF": "Determination of Retained B2O3Content in ZrB2-30 vol_ SiC Oxide Scales.pdf",
  "Text": "['Determination of Retained B2O3 Content in ZrB2-30 vol% SiC Oxide Scales  Kathleen Shugart,*,†,1 Siying Liu, Forrest Craven, and Elizabeth Opila*  Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22904  The  composition of  the borosilicate glass  layer  formed during  oxidation  of ZrB2-30 the extent of B2O3 retention in the oxide during high-temperature oxidation. Oxidation was conducted in stagnant air at 1300°C, 1400°C, 1500°C for  vol% SiC was  determined  to  elucidate  and  times  between  100  and  221 min. Specimens were characterized using mass change and  scanning electron microscopy. After oxidation,  the borosilicate  glass  layer was dissolved from the specimens  sequentially with  deionized H2O and HF acid, using inductively coupled plasma optical emission spectrometry.  to analyze the glass composition  It was  found that  the average B2O3 content to 47 mol%. Retained B2O3 the glass decreased with increasing temperature,  in the glass  scale  ranged  from 23  content  in  the  bulk of  con firming increased volatility with temperature. Boron depth pro files were also obtained in the near  surface region using X-ray  photoelectron spectroscopy and energy dispersive spectroscopy.  The measured B concentrations were  used  to  estimate  the  B2O3 borosilicate  concentration profile and B diffusion coefficients  in the  glass.  Implications  for  the ZrB2-SiC oxidation  process are discussed.  I.  Introduction  O XIDATION  resistance  is  of  great  importance  to Ultra High-Temperature Ceramics (UHTCs), as they are pro posed as Thermal Protection Systems  for hypersonic flight  vehicles. The vehicle leading edges may experience temperatures as high as 2000°C in oxidizing atmosphere.1,2 Studies  on  the  oxidation  characteristics of the UHTC ZrB2 the mid-1960s3-7 with the intent improving the oxidation resistance of this class of materials.  have  been  conducted  since  of  The addition of SiC to ZrB2 rial’s oxidation resistance at  is known to increase the mate higher  temperatures  through  the  formation of a protective borosilicate glass  layer.2,4,8-12  A detailed understanding of  the phases  formed upon oxida tion,  their  composition,  and  the  kinetics  of  formation  are  required as a function of  temperature to enable accurate life  prediction.  The phases present  (in mole  fraction)  after oxidation of  ZrB2-30 vol% SiC are given by the following reaction:  0:61ZrB2 þ 0:39 SiC þ 2:11O2 (g) ¼ 0:61ZrO2 þ 0:39 SiO2 (s/l) þ 0:39 CO(g) þ 0:61B2O3 (l/g)  (1)  where 39 mol% SiC corresponds to the 30 vol% used in this  study. The oxides  form a layered structure, with an outer most  layer of borosilicate glass. A typical example is  shown  in  Fig. 1.  Previous  studies  have  suggested  that ZrB2-SiC  recession rates are controlled by diffusion of oxygen through  the outer glass phase to the base material/oxide interface where the oxidation reactions occur.4,13,14 The B2O3 to SiO2 ratio predicted by Equation 1 is insufficient to define the  glass  composition due  to significant differences  in the vapor  pressures of the two oxides. The vapor pressure of B2O3 at 1500°C is 233 Pa while SiO2 is 3 9 10 suggesting selective volatilization of B2O3 and enrichment of SiO2 in the borosilicate layer relative to the amounts predicted by Equa \\x004 Pa,15,16  tion [1]. Possible B2O3 volatilization reactions include:  B2O3 ð1Þ ¼ B2O3 (g)  (2)  1  2  B2O3 ð1Þ þ 1 2  H2O(g) ¼ HBO2 (g)  (3)  Small changes in B2O3 concentration are shown to have a dramatic effect on oxygen transport in borosilicate glasses. At 1000°C, as little as 1-3 mol % B2O3 in SiO2 on a Si substrate can increase the oxidation rate of Si by two to three orders of magnitude.17 Thus, an understanding of  the B2O3 to predicting  content  in the borosilicate glass  layer  is critical  the oxidation rate of ZrB2-SiC. Oxygen transport in borosilicate glass may occur by per meation or network exchange. Permeation occurs by trans port of molecular oxygen through interstitial  spaces  in the  amorphous borosilicate structure. Network exchange,  similar  to  lattice  diffusion  in  crystalline  phases may  also  occur,  whereby the oxygen jumps from one network site to the next  in the borosilicate glass. Oxygen transport  through pure sil ica has been studied in some detail. Permeation of oxygen in  thermally grown SiO2 on silicon has been shown by Deal and Grove.18 In addition, Ogbuji and Opila19 have found  that  the  enthalpy for  the oxidation of SiC is  essentially the  same  as  that  for Si,  indicating  that  the  controlling mecha nism of oxidation is  the  same  for both materials. Network  diffusion is shown to be much slower than permeation during  oxidation of  silica-forming materials and in silicate glasses, increases.20-22 exchange was shown to  though  the  difference  lessens  as  temperature  A transition  from permeation occur above 1200°C et al.,23,24 while for vitreous silica, Kalen et al. found that both network and interstitial mechanism were in operation.25  to  for  single-crystal  SiC  by  Zheng  However,  few studies have been conducted at  temperatures  as high as those studied here, nor on borosilicate glasses.  The diffusion of B through borosilicate glass  is dependent  on temperature and glass composition. According to Kawaglattice.26  ishi  et al., B diffuses  through Si  sites of  the SiO2 Diffusion rates are found to be on the order of 10 \\x0014 cm2/s, at 10 temperatures ranging between 900°C-1400°C, with low concentrations of B (<15 mol%).26-29 Neither high temperatures nor high B concentrations, have been studied  \\x0020-  for B diffusion in borosilicate glass.  While the specific oxidation rates of ZrB2 and ZrB2-SiC have been measured and semiquantitatively modeled,5-7,14,30-34 a detailed and quantitative understanding of the presence/con centration of B2O3 in the glass layer formed upon oxidation of ZrB2-SiC is currently lacking due to the low sensitivity of most  M. Cinibulk—contributing editor  Manuscript No. 35202. Received June 23, 2014; approved August 27, 2014.  *Member, The American Ceramic Society. 1UES, Inc.; Dayton, OH 45432.  †  Author to whom correspondence should be addressed. e-mail: kns9a@virginia.edu  287  J. Am. Ceram. Soc., 98 [1] 287-295 (2015)  DOI: 10.1111/jace.13236  © 2014 The American Ceramic Society  Journal  \\x0c', 'characterization  techniques  to  light  elements.12,15  Limited  information in the literature supports the presence of B in the  oxide scale and surface depletion. Tripp et al. showed the pres ence of B in the glass phase using Electron Microprobe Analy1300°C and 1400°C.4  sis  (EMPA)  on material  oxidized  at  Karlsdottir et al. discuss the flow of B2O3 to the surface of the oxide, which they analyzed using EPMA and Cathodoluminscence (CL).35,36 They also stated that the surface was depleted  in B2O3 due to its high vapor pressure. This depletion has been shown at the oxide surface by Rezaie et al. using Secondary Ion Mass Spectrometry (SIMS).37 Previous work showed the  formation of borosilicate glass at the very initial stages of oxide  formation using X-ray Photoelectron Spectroscopy though the exact composition was not quantified.38  (XPS),  The  aim of  this  research  is  to  confirm the  presence  of  B2O3 in established oxide scales, in the glass layer formed during oxidation at temperatures of 1300°C-1500°C, and to characterize the compositional gradi to quantify its concentration  ents  formed due to B2O3 volatility. This to the development of a more quantitative description of  information is criti cal  the  oxidation  kinetics  and may  lead  to  approaches  for  improving the currently poor oxidation resistance of ZrB2-30 vol% SiC.  II.  Experimental Procedures  ZrB2-30 vol% SiC bars were provided by Dr. Fahrenholtz (Missouri University of Science and Technology).30,39,40 The  specimens were fabricated using attrition milled powders then hot-pressed. WC contamination (~2 wt%) was observed due to the attrition milling, which used WC  which were  milling media in polyethylene jars. Specimens were cut  from  a  series of billets using an automated surface 40 mm 9 4 mm 9 3 mm and were  grinder  into  bars  of  finished  using  1200 grit diamond abrasive. Bars were sectioned using a low speed diamond saw into coupons approximately 6.2 mm in  length.  Following  sectioning,  each  specimen was  cleaned  ultrasonically  using  detergent  in  deionized  (DI) H2O, DI H2O without detergent, acetone and finally ethanol for 2 min each. Prior to oxidation, the surface area and mass of each  specimen were measured.  Oxidation tests were  conducted in a box furnace  (Rapid Temp, CM Furnaces Inc., Bloomfield, NJ) with molybdenum  disilicide  heating  elements  under  stagnant  air  conditions.  Each coupon was placed on an arc of yttria-stabilized zirco nia  (YSZ, Ortech Inc., Sacramento, CA)  tube  to facilitate  specimen placement and removal  in the furnace, and to limit  contamination from the furnace itself, as previously described.31,38 The arc geometry was chosen to limit contact  between the  specimen and the YSZ tube. Specimen place ment  in the furnace was conducted after the furnace reached  and maintained the desired temperature. Opening the furnace  door dropped the operating temperature approximately 200°C, but the temperature climbed back to within 10°C of  the set point  in 10 s or  less,  leveling out at  the set  tempera ture within 5 min. Specimens were removed from the furnace  while  still  at  temperature,  allowed  to  cool  approximately  15 min after oxidation, and then weighed and measured. The  YSZ boats were also weighed before and after exposure and  negligible mass change was observed.  A baseline  oxidation  test was  conducted  at  1500°C for  100 min. Parabolic oxidation kinetics were assumed based on prior work31,38  in which mass  change  follows  the  relation ship:  Dm  SA  ¼  p ffiffiffiffiffiffiffi  kp t  (4)  where  t  is  the  exposure  time, Dm is  the  oxidation mass  change, SA is  the measured surface area, and kp is sured parabolic oxidation rate constant. The exposure times 1500°C temperatures of  the mea necessary  to  create  the  same mass  gain  as  the  100 min specimen were 1300°C and 1400°C by mg2 cm4/h,  determined  for  using  kp on  values of 5.19 and 4.05 prior work.31 Exposure 1300°C and exposed at these time  respectively,  based  times of 1400°C,  221  and 128 min were  calculated for  respectively. Specimens were  and temperature  conditions  to provide  three specimens with  approximately  the  same mass  gain (and approximately  the  same oxide  thickness), but different oxidation temperatures. Specimens were also exposed for 221 min at 1400°C and 1500°C to provide a series of specimens exposed for the same time at 1300°C, 1400°C, and 1500°C. Three trials of each of these tests were conducted to evaluate variability of the  results. One bar of material was used for  each of  the  three  trials for consistency.  Following mass change measurements, each specimen was  placed in a test tube filled with 15 mL DI H2O maintained at 35°C for 24 h to remove water-soluble B2O3. The specimen was then removed, dried via evaporation at room tem perature and weighed. The DI H2O solution was retained for analysis. The specimen was then placed in another test tube also filled with 15 mL DI H2O maintained at 35°C for 24 h. The second leaching was conducted to confirm that the bulk  of the water-soluble B2O3 was removed during the first H2O soak.41 Again, the specimens were removed and dried and  the DI H2O solution was retained for analysis. After the second leaching bath, each specimen was placed in an HF acid solution (48%-51% HF in H2O) at 35°C for 24 h to dissolve remaining SiO2. In the first trial, the specimens were placed in 3 mL of HF solution  for  24 h  and  then  diluted with  12 mL of DI H2O. mens were soaked in a solution of 12 mL of DI H2O plus 3 mL of HF solution at 35°C for 24 h. The specimens were  In the second and third trials,  the speci removed and dried and the HF solutions were retained for analysis. A very minimal, ~0.05 mL, amount of solution was lost at each removal of the specimen and was unavoidable.  The three solutions  for all  specimens were analyzed for B  and/or Si using  Inductively Coupled Plasma Optical Emis sion  Spectrometry  (ICP-OES, Thermo-Scientific, Waltham,  MA), which has a detection limit on the order of 1 ppb for  most  elements. The purity of  the DI H2O gives mental detection limit of 3 ppb for Si. Precise composition  an experi measurements were made  by  comparison  of  the  emission  spectra from each solution to prepared standards of known  concentrations  and  normalization with  an  internal  yttrium  standard. The standards were prepared using solutions of B  and Si with concentration of 1, 10, and 100 ppm. B and Si  quantities were determined in the H2O solutions, but only Si quantities were determined in the HF solutions, since ZrB2 in HF solutions, as observed by Scanning Electron  dissolves  Microscopy  (SEM,  6700F,  JEOL,  Tokyo,  Japan)  and  described  later. The  total mass  of  all  oxidation  products  Fig. 1. Cross-sectional view of ZrB2-30 1500°C for 100 min in stagnant air.  vol% SiC  oxidized  at  288  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '[ZrO2(s), B2O3(l,g), SiO2(s,l), and CO(g)] was adding the observed mass gained from oxidation and the cal calculated by  culated vaporized mass of both B2O3(g) and CO(g) (assuming full oxidation of the C.) The vaporized mass was  together  calculated using the stoichiometric relation from Equation [1]  and  the  total  amount  of  Si measured  from ICP-OES,  as  described in the  results  section. Finally,  each specimen was  carbon coated with the Precision Etching and Coating Sys tem (PECS),  then examined in the SEM and characterized  with  Energy  Dispersive  Spectroscopy  (EDS,  Princeton  Gamma-Tech Inc., Princeton, NJ)  to ensure the borosilicate  glass layer was completely removed.  Chemical analysis by XPS (PhiVersa Probe XPS spectrom eter, Chanhassan, MN) was performed on two additional specimens oxidized at 1500°C for 10 and 100 min to deter mine  surface B and  Si  concentrations  in  the  glass  layer.  Depth  profiling was  conducted  by Ar  sputtering  at  3 kV  energy on a nonspinning specimen. The  compositional data  were collected on the surface and at depths of approximately  50, 100, and 150 nm into the scale. XPS was performed using  monochromated Al Ka radiation of 1486.6 eV, with a bandpass energy of 11.75 eV and an increment of 0.1 eV/step.  The resolution on Ag3d5 is 0.6 eV. The analyzed area each specimen was about 0.5 mm2 and the  for  scans were per formed  under  identical  conditions. Curve  fitting was  done  using the Gaussian-Lorentzian peak and Shirley background  algorithm of the Phi data massaging program MultiPak. EDS (Oxford Instruments Aztec X-MaxN 150, Concord,  MA) was performed on the surface of a specimen oxidized at 1500°C for 100 min using a series of electron beam accelerat ing voltages  (5, 10, and 20 kV) at a constant working dis tance  of  15 mm.  The  change  in  B  composition  with  accelerating voltage and simulated sampling depth was used  to provide additional  information about  the B concentration  profile  in the borosilicate  glass,  as described further  in the  results section.  III.  Results  (1)  B2O3 Concentration in Borosilicate Glass  SEM and EDS results for two of  the specimens after borosil icate  glass  digestion  are  shown  in Fig. 2.  SEM was  per formed  using  a  beam voltage  of  5 kV  unless  otherwise  specified to increase  sensitivity  to elements of  light  atomic  mass during EDS. Comparing the SEM and EDS results of  glass dissolution trials,  it  can been seen that  trial  1  (3 mL  HF, 24 h)  removed all of  the glass and ZrB2 leaving only SiC. Trial 2 used more dilute HF acid to  from the  sur face,  dissolve the glass  layer,  leaving some ZrO2 grains and miniresidual glass on the specimen surface, which were not  mal  found in trial 1. These  results  indicate  that  the assumption  that all  the glass was  removed in the water and HF dissolu tion procedure  is  reasonable,  ensuring the  ICP-OES results  provide an accurate average  composition of  the borosilicate  glass  formed during oxidation. Note that SiC is not  soluble  in H2O or HF as confirmed by SEM (Fig. 2) so that Si content in the solutions must be attributed to SiO2. The specific mass change from the three oxidation trials  and the B2O3 “Mass Change-Oxidation” is the measured change in mass of  leaching process are  reported in Table I. The  the specimen after oxidation under the stated conditions. The  “Mass Change-Leaching”  is  the mass of material  removed  after both of  the H2O soaks. This mass was determined by subtracting the specimen mass after the H2O baths from the expected to be composed of  mass  after oxidation and was  mostly B2O3 with trace amounts of SiO2. The B2O3 retained in the glass layer after oxidation was calculated assuming:  (1)  an  equal  rate  of  consumption  of  the ZrB2 and SiC in the base material during oxidation; the water leaching steps removed all of the B from the layer;41 and (3) all of  (2)  glass  the Si  from the glass  layer was  removed  in  the H2O and HF acid  soaks. The  validity  of  these assumptions is discussed later. For this calculation,  the  concentrations of B and Si  in the glass  layers of each speci men were determined by ICP in mg/L and then multiplied by  0.015 L,  the volume of each solution. This yielded the mass  in mg of dissolved B and Si, which was then easily converted  to moles. Since the SiO2 does not volatilize in this ture range, multiplication of the amount of Si in moles by  tempera the  stoichiometric  ratio  of B to  Si  expected  from Equa tion (1) yields the total amount of B expected if none volatil izes  (see  Appendix  A  for  an  example  calculation.)  Comparison of  the B concentration found from ICP-OES to  the  total  amount oxidized according  to the  above  calcula tions  gives  the  percent  of  retained B. Table II  provides  a  summary of these calculations. Percent B2O3 given by 100%-% retained.  volatilized  is  Figure 3  is  a plot of  the  total mass of oxides  generated  versus  temperature  for  specimens which were oxidized with  the  goal  of  attaining  the  same  amount  of  oxidation. The  total mass of generated oxides includes the mass gained dur ing oxidation and the calculated mass of B2O3(g) and CO(g) which have formed and vaporized. The similar mass results  for the three temperatures and three trials show that  the cho sen times do provide similar quantities of oxidation, allowing  comparison between these  tests. Figure 4  is  a plot of  total  mass of  generated oxides  versus  temperature  for  specimens  all oxidized for  221 min,  showing with 80% confidence  an  increase in the amount of oxide formed with increasing tem perature.  XPS  results  of  oxidized  specimen  surfaces  are  given  in  Fig. 5. For both 10 and 100 min exposures,  less  than 1% B  was evident on the surface, but after sputtering 50 nm,  the B  content  became measureable  and  steadily  increased  with  depth to 150 nm into the  surface,  the deepest  sputter depth  (a)  (b)  Fig. 2.  SEM/EDS results for ZrB2-30 vol% SiC oxidized at 1500°C for 100 min in stagnant air after two DI H2O soaks and dissolution 35°C for in (a) 3 mL of HF solution at 24 h (b) 3 mL of HF solution diluted with 12 mL of DI H2O at 35°C for 24 h, the area with the most products left behind after dissolution.  showing  January 2015  Determination of Retained B2O3  289  \\x0c', 'analyzed. The B content measured by XPS at depths greater  than  50 nm after  100 min  oxidation was  about  5% lower  than  it was  after  10 min  oxidation, while  the  Si  content  increased 5%. The O and Zr values were equivalent and con stant for the two scales.  A boron concentration profile was  estimated using EDS 1500°C for  results  obtained  from a  specimen  oxidized  at  100 min,  at  a  series  of  accelerating  voltages  (5,  10,  and  Table I.  Mass Change Data for ZrB2-30 vol% SiC Oxidized Under Stagnant Air Conditions After Oxidation and After H2O Soak (Leaching)  Condition  Trial #  Mass change - Oxidation (+mg/cm2)  Average (+mg/cm2)  Mass change - Leaching (-mg/cm2)  Average (-mg/cm2)  1500°C 100 min  1  2.81  2.90 \\x06 0.26  0.70  0.91 \\x06 0.28  2  3.12  1.23  3  2.70  0.80  1400°C 128 min  1  3.85  4.55 \\x06 0.68  1.13  2.07 \\x06 1.05  2  5.21  3.20  3  4.60  1.88  1300°C 221 min  1  3.46  4.61 \\x06 1.06  2.39  2.38 \\x06 0.37  2  5.52  2.75  3  4.86  2.00  1400°C 221 min  1  6.28  7.10 \\x06 1.04  4.29  3.94 \\x06 0.45  2  8.27  4.11  3  6.76  3.43  1500°C 221 min  1  4.38  4.91 \\x06 0.49  1.48  2.02 \\x06 0.74  2  5.00  1.71  3  5.36  2.87  Table II.  Summary of B Retained in the Borosilicate Scale After Oxidation of ZrB2-30 vol% SiC Under Stagnant Air Conditions. A High Percentage of the B does not Volatilize under the Conditions Tested, Leading to a High B2O3 Concentration in the Borosilicate Glass  Conditions  Trial #  B2O3 content  of glass (mol%)  Average B2O3  content  (mol%)  % of Total B retained  in glass (mol%)  Average B  retained (mol%)  “Constant”  Oxide thickness  1500°C 100 min  1  27.3  27.3 \\x06 4.1  23.9  22.6 \\x06 4.7  2  29.2  26.4  3  21.3  17.3  1400°C 128 min  1  31.9  36.8 \\x06 5.7  28.5  37.3 \\x06 10.1  2  43.1  48.3  3  35.5  35.1  “Constant”  time  1300°C 221 min  1  43.2  42.9 \\x06 1.6  43.3  46.2 \\x06 4.0  2  44.3  50.8  3  41.1  44.6  1400°C 221 min  1  34.2  36.9 \\x06 2.4  37.3  38.8 \\x06 1.9  2  39.0  40.9  3  37.4  38.1  1500°C 221 min  1  27.3  28.2 \\x06 2.0  24  25.1 \\x06 2.5  2  26.8  23.4  3  30.4  28  Fig. 3.  Total mass of generated oxides versus temperature for ZrB230 vol% SiC specimens oxidized to generate the same oxide thickness. 1300°C specimens oxidized for 221 min, 1400°C specimens oxidized for 128 min, and 1500°C specimens oxidized for 100 min.  Fig. 4.  Total mass of generated oxides versus temperature for ZrB230 vol% SiC specimens oxidized for 221 min.  290  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '20 kV.) The Casino simulation program was used to deter mine the depth of voltage.42 These  the sampling volume for each accelerating  values  are  given in Table III. A series of  layers were  simulated with  the Casino  program using  the  compositional data determined both from XPS and EDS and the density data from Bruckner et al.41 The sampling volume  for  the EDS was determined based on the  layered borosili cate structure. The first  layer of  the final structure was set  to  be  50-nm thick  with  a  composition  of  6.1 mol% B2O3 from averaging the XPS results for the sur (remainder SiO2) face and 50 nm into the borosilicate. This composition has a density of 2.21 g/cm3. The next 100 nm was assumed to have  a composition of 8.8% B2O3 (remainder SiO2) ing XPS results acquired at 50 and 150 nm depth, with a 2.18 g/cm3. The next  from averag corresponding density of  three  layers  were  assumed  to  have  a  composition  of  13.0,  18.7,  and  31.0 mol% B2O3 (remainder SiO2) positions were determined from mol% B and Si measured by  respectively. These  com EDS using  acceleration voltages of  5,  10,  and 20 kV. The  last  three  layer  thicknesses were adjusted using an iterative  process  to determine  the  average  sampling depth  for  each  accelerating voltage. Table IV provides the layer information  used to generate  the Casino simulations. The  three Casino  simulations  of  the  sampling  volume  at  5,  10,  and  20 kV  based on final  iterations of  the layered structures were used  to provide B concentration and depth values. An example  simulation for  5 kV is given in Fig. 6. The final depth for  each  acceleration  voltage was  determined  using  plots  of  depth versus normalized hits  (electrons hitting atoms  in the  material), as given in Fig. 7 for 5 kV. Estimated B concen trations  from the EDS results  are plotted along with XPS  data in Fig. 8.  (2)  Calculated B Concentration Gradient  The borosilicate glass  layer  formed after oxidation of ZrB230 vol% SiC was approximated as a binary, semi-infinite sys tem, and the error function solution to Fick’s 2nd Law was used to calculate the B concentration gradient:43  CB ¼ C1 erfðgÞ  (5)  g ¼  x  4Dt  (6)  erfðgÞ ¼ 2ffiffiffi p  p  Z  g  0  e  \\x00k2  dk  (7)  where CB borosilicate  is  the  concentration of B at distance x into the  glass  from the  gas  interface, C∞ is  the  equilib rium concentration of B at  the oxide/base material  interface,  D is  the diffusion coefficient of B in the borosilicate  glass,  and t is time. The following boundary conditions were used: BC1: CB = 0 at x = 0 BC2: C∞ = 0.289mole fraction (fixed by ZrB2-30 vol% SiC) at x = 20.2 lm These boundary conditions are reasonable as XPS shows a very low B concentration (<1%) on the  stoichiometry of  surface, and O2(g) diffuses through the borosilicate and forms new oxide rapidly  in comparison to B diffusion outward. The concentration of  B at  the borosilicate glass/base material  interface is assumed  constant. Choosing  a fixed time,  the B diffusion coefficient  was determined using:  Fig. 5.  XPS results for ZrB2-30 vol% SiC oxidized at 1500°C for 10 (open symbols) and 100 (closed symbols) min in stagnant air.  Table III.  Concentration Results from EDS Performed at  Varying Accelerating Voltages, Showing Increased B  Concentration Deeper in the Borosilicate Scale  Accelerating  voltage (kV)  B concentration  (at%)  B2O3 concentration  (mol%)  Simulated average  sampling depth (nm)  5  7.6  13.0  220  10  11.0  18.7  770  20  18.5  31.0  2900  January 2015  Determination of Retained B2O3  291  \\x0c', 'DB ¼ ½ð1 \\x00 Cavg t  C1  Þ  ffiffi  p  p  xl  2  \\x8a2  (8)  Solutions were obtained for a specimen oxidized at 1500°C after 100 min, which forms a borosilicate layer 20.2 lm thick  (BC2),  since  ICP-OES, XPS, and EDS compositional  infor mation are available  for  this oxidation condition. The aver age  B  concentration  in  the  borosilicate  glass  after  this  oxidation treatment given by ICP-OES was 15.4 mol%, and  the concentration at  the oxide/base material  interface is assu med to be  equal  to that  given by  equilibrium, 28.9 mol% found to be ~1.2 9 represents an average  (BC2). The B diffusion coefficient was \\x0010 cm2/s. This diffusion coefficient 10 B diffusion coefficient in borosilicate glass resulting from ZrB2-30 vol% SiC oxidation at 1500°C for 100 min. B diffusion coefficients were also calculated which provided best fits \\x0013 cm2/s to the XPS data and the EDS data and were 1 9 10 \\x0012 cm2/s, respectively. The DB values will be disand 5 9 10 cussed below. Resulting B concentration profiles are shown  in Fig. 8.  IV.  Discussion  (1)  B2O3 Concentration in Borosilicate Glass  Assuming  a  linear  concentration gradient  from 0.289 mole  fraction B (0.61 mole  fraction B2O3) at diboride interface (which corresponds to ZrB2-30 vol% SiC) interface, the average  the glass/zirconium  to  zero  at  the  gas/borosilicate  glass  B2O3 mole fraction is 0.305 which is results obtained at 1400°C and 1500°C,  intermediate  to  the  shown in Table II.  This  suggests  that  the  ICP-OES results are  reasonable. The  ICP-OES results 1500°C relative  confirm a 1300°C,  higher  depletion  of  B2O3 expected B2O3 gain behavior of the speci at  to  consistent with  volatility. Comparing  the mass  mens (Table I) to the total mass of oxides generated in Fig. 3  corroborates this B2O3 volatility.  (2)  Calculated B Concentration Gradient  The B diffusion coefficient  is expected to vary with borosili cate  composition  and will  therefore  vary with  position  in  the oxide  scale. The  coefficient determined for  the  average 1500°C  composition  from  ICP-OES after \\x0010 cm2/s. 1.2 9 10 the surface  oxidation  at  for  100 min  was  The  estimated  EDS  data, which  is  nearer  to  and  thus  lower  in B  content, was  best  fit with  a  boron  diffusion  coefficient  of  Table IV. Layer Information Used in Casino Simulations for EDS Penetration Depth Determinations of Borosilicate Glass Formation on ZrB2-30 vol% SiC After Oxidation at 1500°C for 100 min in Stagnant Air in Standard Box Furnace  Layer  Thickness  (nm)  Total depth  (nm)  Measured B  concentration (at%)  Composition  (mol% B2O3)  Assumed density (g/cm3)  Source of B data  1  50  50  0.1, 3.8  6.1  2.21  XPS at surface and 50 nm  2  100  150  3.8, 5.4  8.8  2.18  XPS at 50 nm and 150 nm  3  70  220  7.6  13.0  2.15  EDS, 5 kV  4  550  770  11  18.7  2.12  EDS, 10 kV  5  2130  2900  18.5  31.0  2.06  EDS, 20 kV  Fig. 6.  Casino  simulation  showing EDS  sampling  volume  for  an  accelerating voltage of 5 kV in borosilicate glass formed on ZrB2-30 vol% SiC after oxidation at 1500°C for 100 min in stagnant air in  standard box furnace.  Fig. 7.  Casino simulation of depth versus normalized hits generated  for 5 kV in the simulated borosilicate glass.  Fig. 8. XPS and EDS results 5 lm of 1500°C  showing B concentration in the first  the  borosilicate  glass  scale  for  a  specimen  oxidized  at  for  100 min  in  stagnant  air. Dashed  line  indicates  B  concentration  profile  estimated  using B diffusion  coefficient  from  best fit  to XPS results. Dotted line indicates B concentration profile  estimated using B diffusion coefficient  from best fit  to EDS results.  Solid line indicates B concentration profile estimated using calculated  B diffusion coefficient from ICP-OES results.  292  Journal of the American Ceramic Society—Shugart et al.  Vol. 98, No. 1  \\x0c', '\\x0012 cm2/s. The compositional data from the XPS rep5 9 10 the lowest B concentration due to proximity to the  resents  surface and was best fit with a boron diffusion coefficient of \\x0013 cm2/s. All of 1 9 10 than previous measured values for B diffusion through boro\\x0014 cm2/s when extraposilicate, which are on the order of 10 lated to 1500°C.27,29,44 The high values of the B diffusion  these diffusion coefficients are higher  coefficient estimated in this work are reasonable since boro1500°C as silicate glass with 27 mol% B2O3 is liquid at shown in the SiO2-B2O3 phase diagram (Fig. 9).45 The \\x0013 cm2/s, mated B diffusion coefficient of 1 9 10 obtained from best fit to the XPS data is reasonable as these data were  esti taken near the surface where less B is present. The glass here  will be more viscous and lower diffusion rates are predicted.  Figure 8 compares B concentration profiles based on best fit  for  the  ICP-OES average  composition, XPS data, and EDS  estimates.  Note also that the average B concentration determined by ICP-OES is a lower bound for B concentration. Bruckner41  states that 95% of B2O3 is removed from borosilicate glasses by a warm water leach, however, Adams and Evans46 state  that only 30% is  removed.  It  is not possible to confirm the  extent  of B2O3 ZrB2 dissolves in HF. If B leaching in H2O is incomplete, average B composition in the borosilicate scale would  leaching  in warm water  at  this  time  since  the  be  higher  and  the  calculated B concentration would  increase,  bringing it  in closer agreement with the XPS and EDS gradi ent results in Fig. 8.  A B2O3 concentration gradient proposed in Fig. 10, for the first  in the borosilicate layer  is  time  providing  composi tional  information about  this  layer. The B2O3/SiO2 is assumed to be equal  ratio at  the oxide/substrate  interface  to the  B/Si  ratio  in  the  base material. This  is  justified  by  two  observations. First, XPS analysis  showed that B-O and Si-O  bonds are both observed on the oxidized surface after only  10 s, indicating both oxides are forming simultaneously at the start of oxidation.38 Second, a nearly uniform oxidation  front  is  observed  between  ZrB2/SiC and the oxides, as shown in Fig. 1. The presence of the ZrO2+C layer between  the glass  layer and the base material  is neglected in Fig. 10  and is  assumed to have  little  effect on the B/Si boundary  condition at  the  interior  glass  interface  (K. N. Shugart &  E.  J. Opila,  in preparation). The B2O3 gradient calculated by converting the B concentration gradient  in the plot  was  estimated from ICP-OES in Fig. 8 to oxide. The discontinuborosilicate/ZrB2+SiC interface the average B concentration determined by ICP-OES  ity  at  the  results  from the  use of  to calculate an average diffusion coefficient.  (3)  Implications for Oxidation  It  is confirmed here that considerable B2O3 is retained in the glass layer which has important effects on oxidation, assum ing the rate-limiting step is oxygen diffusion in the borosili1500°C, layer.4,17 At  cate  glass  the  liquidus  curve  lies  at  approximately 95 mol% B2O3, as surface of the glass layer depleted in B2O3 will be composed of a more viscous SiO2 then the interior of the glass layer with higher B2O3 content. Temperature transport will be  seen in Fig. 9. Thus,  the  effects  on  oxygen  very  complex.  Increasing  temperature  depletes the borosilicate glass in B2O3, slowing oxygen transport by compositional effects at the surface, but at the same  time  increases oxygen transport  since  it  is an activated pro cess.  The  apparent  activation  energy  for  oxidation  then  reflects  some combination of B2O3 vaporization, B transport outward due to the driving force induced by the B2O3 concentration gradient, and oxygen transport inward in the non uniform glass structure.  V.  Conclusions  Significant B2O3 scale formed during oxidation of ZrB2-30 vol% SiC at temperatures between 1300°C and 1500°C. Borosilicate composi is  found (27-43 mol%)  in the borosilicate  tions measured using ICP-OES (average value), EDS & XPS  were used to estimate B diffusion coefficients \\x0013 cm2/s to 10 for B2O3-rich and respectively. The variation in  in the borosili cate  ranging  from 10  \\x0010  SiO2-rich portions of B concentration across  the scale,  the oxide  scale  implies variations  in  Tm, DB, and DO through the resulting complex temperature dependence for oxygen perme thickness of  the  scale  and a  ation,  the likely rate controlling mechanism for oxidation of  ZrB2-SiC.  Fig. 9.  B2O3-SiO2 Equilibria Diagrams,  phase  diagram,  from  ACerS  NIST  Phase  figure  11777.  (Reprinted with  permission  of  The American Ceramic Society.)  Fig. 10.  Approximate  B2O3 formed during oxidation of ZrB2-30 vol% specimen oxidized at 1500°C for 100 min  concentration  (mol%)  profile  in  borosilicate  glass  scale  SiC. Values given are for  in  stagnant  air,  using  average  B  concentration  (ICP-OES)  to  determine  the  average  diffusion  coefficient.  A majority  of  the  borosilicate glass is shown to have sufficient B2O3 to be liquid at oxidation temperature.  the  January 2015  Determination of Retained B2O3  293  \\x0c', 'Appendix A  Oxidation reaction: 0.61 ZrB2+0.39 SiC+2.11 O2(g)=0.61 ZrO2+0.39 SiO2(s 0.39 CO(g)+0.61 B2O3(l/g) Example calculation for: 1500° C, 100 minutes, Molar mass for B2O3= 69.62 g/mole Molar mass for CO= 28.00 g/mole  /l)+  trial 1  Si  to B ratio in starting material:  B  Si  ¼ 0:61mol \\x03 2 0:39mol  ¼ 3:13  [A1]  Total Si  in SiO2 measured via ICP-OES:  2:97mg/L þ 0:95mg/L þ 86:05mh/L ffi 89:97mg/L  [A2]  (H2O soak 1 + H2O soak 2 +HF soak= total Si) Measured Si (mol):  89:97mg/L \\x03  0:015L  28:09g/mol  \\x03  1g  1000mg  ¼ 4:80 \\x03 10 \\x005mol  [A3]  Total B (mol)  consumed  during  oxidation,  predicted  from  [A1] and [A3]:  4:80 \\x03 10 \\x005mol \\x03 3:13 ¼ 1:50 \\x03 10 \\x004mol  [A4]  Total B measured via ICP-OES:  22:03mg/L þ 3:91mg/L ¼ 25:94mg/L  [A5]  (H2O soak 1 + H2O soak 2 = total B) Measured B (mol):  25:94mg/L \\x03  0:015L  10:81g/mol  \\x03  1g  1000mg  ¼ 3:60 \\x03 10 \\x005mol  [A6]  Percentage of B retained ([A6]/[A4]):  3:60 \\x03 10 \\x005mol 1:50 \\x03 10 \\x004mol  ¼ 0:24 ¼ 24%  [A7]  Vaporized B (mol):  1:50 \\x03 10 \\x004mol \\x00 3:60 \\x03 10 \\x005mol ¼ 1:14 \\x03 10 \\x004mol  [A8]  Total B2O3 and CO vaporized:  1:14 \\x03 10 \\x004mol  2 \\x03 1000mg 1g  \\x03 69:62g mol  þ 4:80 \\x03 10 \\x005mol \\x03 28g mol  \\x12  \\x13  ¼ 5:31mg  [A9]  Surface area of specimen: ð6:17mm \\x03 3:99mmÞ \\x03 2 þ ð6:17mm \\x03 2:97mmÞ \\x03 2 þ ð2:97mm \\x03 3:99mmÞ \\x03 2 ¼ 109:59mm2  [A10]  Measured mass after oxidation minus measured mass before  oxidation:  384:95mg \\x00 381:87mg ¼ 3:08mg  [A11]  Total mass of generated oxides  (ZrO2+SiO2+B2O3+CO) per  area:  ð3:08mg þ 5:31mgÞ 109:59mm2  \\x03 100  mm2  cm2  ¼ 7:67mg/cm2  [A12]  Acknowledgments  The authors acknowledge  Joe Hagan and Bohuslava McFarland (University  of Virginia)  for  the ICP-OES results and Dr. Wayne Jennings  (Case-Western  Reserve University)  for  the XPS results. The authors would like to thank Dr.  William Johnson for helpful advice and discussions on calculating the boron  diffusion coefficient. The authors would also like  to acknowledge Eric Neu man  and Dr. William Fahrenholtz  at Missouri University  of  Science  and  Technology for  the ZrB2-30 vol% SiC material. by The National Hypersonic Science Center-Materials  Initial  funding for  this work  was  provided  and  Structures.  References  1A. Paul, D. D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. G. P. Bin ner, , B. Vaidhyanathan, A. Heaton, P. Brown, and W. E. Lee, “UHTC Composites for Hypersonic Applications,” J. Am. Ceram. Soc., 91, 22-28 (2012). 2A. Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby, “Oxidation  of ZrB2-SiC Ceramics Under Atmospheric and Reentry Conditions,” Refract. Appl. Trans., 1, 1-7 (2005). 3H. C. Graham, H. H. Davis, I. A. Kvernes, and W. C. Tripp, “Microstruc tural Features of Oxide Scales Formed on Zirconium Diboride Materials,” Mater. Sci. 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},{
  "_id": 41,
  "PDF": "Effect of carbon nanotube on processing, microstructural, mechanical and ablation behavior of ZrB 2 -20SiC based ultra-high temperature ceramic composites.pdf",
  "Text": "['Carbon 111 (2017) 269e282  Contents lists available at ScienceDirect  Carbon  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c a r b o n  Effect of carbon nanotube on processing, microstructural, mechanical and ablation behavior of ZrB2-20SiC based ultra-high temperature ceramic composites  Ambreen Nisar a, S. Ariharan a, T. Venkateswaran b, N. Sreenivas b, Kantesh Balani a, *  a Department of Materials Science and Engineering, b Vikram Sarabhai Space Centre,  Indian Institute of Technology Kanpur, Kanpur, 208016,  India  Indian Space Research Organization, Trivandrum, 695022,  India  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 15 July 2016 Received in revised form 30 September 2016 Accepted 1 October 2016 Available online 4 October 2016  Herein, zirconium diboride (ZrB2) is reinforced with silicon carbide (SiC) and carbon nanotube (CNT) to provide enhanced structural stability and oxidation protection against extreme thermal (>2400 \\x0e C) and oxidative environments. The ablation resistance of ZrB2-based composites was evaluated using plasma arc-jet with a heat-flux of 2.5 MW/m2 for 30 s, and the decreased oxidation-rate (from 0.77 mm/s to 0.44 mm/s) is attributed to enhanced thermal conductivity (42.3e52.3 W/mK at 1200 \\x0e C) with synergistic reinforcement of SiC and CNT. The increased onset temperature (from 679 \\x0e C to 706 \\x0e C) and decreased enthalpy of oxide formation (from 1.6 to 0.6 kJ/g), insinuates an increase in thermal stability and oxidation resistance with the synergistic addition of both SiC and CNT in ZrB2. The increase in the hardness of ZrB2 in the as-processed composites (up to 1.6 times) as well as after plasma arc jet exposure (up to 2.1 times) with synergistic reinforcement of SiC and CNT has shown to suppress crack-formation and restrict oxidation. The reduction in the analytically evaluated tensile interfacial residual stress indicates enhanced structural integrity of ZrB2-SiC-CNT composites, which is a mandatory requirement for aerospace applications.  © 2016 Elsevier Ltd. All rights reserved.  1.  Introduction  Zirconium diboride (ZrB2), an ultra-high temperature ceramic (UHTC) possesses low density (6.09 g/cc) [1], high melting temperature (Tm ¼ 3245 \\x0eC) [2] and high thermal conductivity (60 W/ mK) [3], which makes it attractive for use in thermal protection system and scramjet engine components for hypersonic vehicles. The advantages of the ZrB2-based UHTCs is not only from their high-temperature stability, but also from their high thermal conductivity at elevated temperatures [3], which ensure its utilization as sharp leading edges for re-entry space vehicles. The superior high temperature stability is primarily due to the strong covalent bonding in ZrB2 structure [3], however, strong bonding and low diffusion coefficient make the processing of ZrB2 a challenge. Reinforcements are often added to ZrB2 to improve densification, mechanical (hardness, elastic modulus, and fracture toughness), thermal and oxidative properties [4e6]. Several investigations  * Corresponding author. E-mail address: kbalani@iitk.ac.in (K. Balani).  http://dx.doi.org/10.1016/j.carbon.2016.10.002 0008-6223/© 2016 Elsevier Ltd. All rights reserved.  indicated that SiC reinforcement not only improves flexural strength, fracture toughness but also improves oxidation properties by the formation of passive SiO2 layer which protects ZrB2 at elevated temperature [6e11]. It has been established that ZrB220SiC (vol%) composites is optimal for high-temperature thermomechanical applications [9,12]. On one hand, where the thermal resistance of SiC reinforcement is well established, reinforcement of thermally conductive CNT (~3000 W/mK) [13,14] may provide enhanced resistance to thermal damage of ZrB2-composites [6]. Recently, carbon nanotube (CNT), graphite and carbon fibers have also been used as nano-scale filler in the ZrB2 ceramics [6,15,16], resulting in nearly complete densification [6,16]. Tian et al. [15] investigated the effect of CNT (2 wt%) on the properties of ZrB2SiC ceramics, showing that the fracture toughness increased about from 4 MPam1/2 4.6 MPam1/2), 15% (i.e. to but hardness (15.8 GPae15.5 GPa), and thermal conductivity (at 400 \\x0eC, 72.5 W/ mK to 75.4 W/mK) did not vary significantly. Plasma-arc jet testing closely approximates the intense aerothermal conditions [17] which typically exceeds 2000 \\x0eC during the test and provides high shear stresses (due to the flux of 1e3 MW/ m2) that can influence the microstructural evolution (enough to  \\x0c', '270  A . Nisar et al.  / Carbon 111 (2017) 269e282  melt most UHTC oxides) leading to mechanical failure. While improving mechanical properties of ZrB2 is the crucial step in making it a viable material, it is not sufficient, if the oxidation properties are not improved upon. While evaluating the oxidation behavior of these UHTCs, it is necessary to evaluate the materials performance by simulating environmental conditions experienced during hypersonic flight. Till date, the oxidation mechanism of only monolithic ZrB2 is available [18], but, ZrB2 reinforced with 20 vol% SiC (ZrB2-20SiC) is established as the baseline material in literature, so a comprehensive understanding of the effect of microstructural evolution on thermo-mechanical performance (of UHTC) is required for their use in hypersonic and propulsion applications. In accordance with the recent study on TaC-based composites [19,20], the synergy of SiC and CNT reinforcement has been extended to ZrB2-based composites in the current work. The scope of this work is to compare the effect of CNT on the oxidation behavior of ZrB2-20SiC under plasma arc jet exposure. In agreement with the previous study in literature [6,15], where up to 6 vol% of CNT is added in ZrB2 and ZrB2-20SiC, these authors have observed an increasing trend of fracture toughness even with the maximum CNT reinforcement of 6 vol % (i.e. 3.53 MPam1/2 in ZrB2, and 4.6 MPam1/2 in ZrB2-20SiC composite). Thus, the aspect of agglomeration may not have occurred in these composites. On that basis, 10 vol% (which corresponds to 4.1 wt %) of CNT was reinforced in order to improve densification behavior, mechanical properties, thermal conductivity, and oxidation resistance of ZrB2-20SiC. Enhanced content of CNT may protect the base ZrB2 material via sacrificial oxidation at high temperatures. Thus, the consequence of CNT addition on the high temperature thermal conductivity is correlated with the oxidation behavior of the ZrB2-based composites. The composition of the surface layer and variation in the thickness of oxide scale (after plasma arc jet exposure) is also investigated to analyse the oxidation resistance of ZrB2-SiC-CNT UHTC composites. The oxidation mechanism has been proposed according to the experimental results and thermodynamic considerations, based on which sample ranking has been made against the performance upon plasma arc jet exposure. Furthermore, the thermogravimetric analysis is carried out at 1500 \\x0eC (for 30 min) in an oxygen environment, to complement the results obtained from plasma arc jet exposure. Through such investigation, the purpose is to correlate the effect of microstructural evolution on the structural integrity (with the synergistic addition of SiC and CNT) during processing, which governs the oxidative performance of ZrB2-based UHTC against plasma arc jet testing.  2. Materials and method  2.1. Materials processing  Commercial powders of ZrB2 (H.C. Starck, Germany, 99.9% pure, particle size < 2 mm, Fig. 1a), SiC (H.C. Starck, Germany, 99.9% pure, particle size < 1 mm, Fig. 1b) and multi-walled carbon nanotubes, (CNT, Nanostructured and Amorphous Materials Inc., TX, USA, 94% pure, outer diameter of 50 nm, inner diameter of 30 nm, and 1e2 mm long, Fig. 1c) were used as the starting materials. The powder mixtures: ZrB2 mixed with 20 vol% SiC (referred to as Z20S), ZrB2 with 10 vol % CNT (referred to as Z10C), and ZrB2 with 20 vol% SiC and 10 vol% CNT (referred to as Z20S10C) were dry ball milled with ball to powder ratio of 5:2 for 8 min at 500 rpm using tungsten carbide jar and tungsten carbide balls. Bright-field TEM micrograph of Z20S10C (Fig. 1d) illustrates the distribution of secondary phase particles (SiC as well as CNT) in ZrB2 matrix and the retention of CNT after ball-milling.  Fig. 1. TEM micrographs of staring powder (a) ZrB2 (b) SiC (c) CNT and (d) Z20S10C composite powder.  The ball-milled composite powders were processed via spark plasma sintering (SPS, Dr. Sinter, SPS-515S, Japan) for resultant pellet of 15 mm diameter and 2e3 mm thickness using graphite die and punches with heating rate of 150 \\x0eC/min by holding 5 min at each step (1300 \\x0e C (0.4 Tm), 1500 \\x0e C (0.45 Tm) and 1650 \\x0eC (0.5 Tm)) at 40 MPa in argon (Ar) atmosphere. After final stage of holding at final temperature, the sample was allowed to cool naturally in the presence of Ar gas. The theoretical density of these composites were calculated using rule of mixture (ROM), however, the final densities were measured by Archimedes method (using distilled water as immersion medium) on a hydrostatic balance.  Fig. 2. Plot showing the variation of instantaneous densification with time processed in subsequent stages during SPS processing of ZrB2-based composites. (A colour version of this figure can be viewed online.)  \\x0c', 'A . Nisar et al.  / Carbon 111 (2017) 269e282  271  Table 1 Nomenclature, theoretical and experimental densities, and % relative densification of ZrB2-SiC-CNT composites.  Sample ID  ZrB2 Z20S Z10C Z20S10C  Composition  ZrB2 ZrB2þ20 vol%SiC ZrB2þ10 vol%CNT ZrB2þ20 vol%SiCþ10 vol%CNT  rtheo (g/cc)  rexpt (g/cc)  % Densification  6.1 5.5 5.7 5.1  5.7 5.2 5.5 5.0  93.1 95.0 95.9 99.7  flux and operating conditions of the plasma generator with argon calibrated using water cooled Gardon gauges. The linear oxidation rate (R) is calculated using the formula:  R ¼ doxide scale texposure  (1)  doxide  where, is the oxide scale thickness after plasma arc jet exposure and texposure is the exposure time (30 s). Complimentarily, simultaneous thermal analysis (STA 8000, Perkin Elmer, USA) i.e. thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) is used to evince oxidation mechanism and thermal stability. Thermal studies were performed on samples with an initial mass of 50e100 mg from room temperature to 1500 \\x0eC at a rate of 20\\x17C/min in an oxygen atmosphere for 30 min holding at maximum temperature. The accuracy of the equipment is 0.2% with furnace within ±0.5 \\x0eC. Under the thermal stability of similar conditions, the reported value is an average of at least 5 experiments.  2.3.  Thermal conductivity measurements  The thermal conductivity of the investigated samples was measured by laser flash diffusivity technique (Flashline, thermal properties analyzer, Anter Corporation, USA) using pulsed Nd:YAG laser for ~300 ms in Ar atmosphere. The back face temperature rise of the samples (diameter of 14 mm and thickness of 3 mm) was recorded by In-Sb photo detector. The thermal diffusivity, D is calculated using relation (Eq. (2)).  D ¼ 0:13879 L2  t1=2  (2)  where, L is specimen thickness and t1/2 is the half time required for initiation of the pulse for back face temperature to reach half of the maximum rise in temperature. Thermal conductivity, k is calculated using the relation (Eq. (3)) [21]:  k ¼ D r Cp  (3)  where, r is the measured density, Cp is the specific heat (simultaneously measured using laser flash diffusivity technique with an accuracy of ±4%). It may be indicated that the effect of CNT agglomeration is not accounted for during the thermal conductivity measurements. The thermal diffusivity, specific heat, and thermal conductivity of these materials were analyzed under inert atmosphere in order to better understand the influence of porosity and composition. The true thermal diffusivity values are estimated to be within an error of ±1.44% with 95% confidence limit.  2.4.  Phase, microstructural and mechanical characterizations  Phase analysis was carried out using Rich-Seifert, 2000D diffractometer operated at 25 kV and 15 mA (Cu Ka, l ¼ 1.54 Å at a scan speed of 0.5 s/step and a step size of 0.02\\x0e ) in the 2q range from 30\\x0e to 90\\x0e . The quantification of the retained ZrB2 phase after  Fig. 3. (a): XRD pattern of initial composite powders as well as ZrB2-SiC-CNT composite pellets, and (b) Raman spectra showing Dand G-peaks in pristine CNT powder as well as SPS ZrB2-SiC-CNT pellets. (A colour version of this figure can be viewed online.)  2.2. Oxidation test  SPS samples of 15 mm diameter and 3 mm thickness, encased in carbon phenolic guard, were plasma arc-jet exposed under heat flux of 2.5 MW/m2 for 30 s. The heat flux (2.5 MW/m2) is directly measured using water cooled thin foil heat flux transducers (HFT) which has an uncertainty of ±5%. The HFT is placed at a desired axial location from nozzle exit and power and gas flow to plasma generator is adjusted to get this heat flux. To monitor the back face temperature, a K-type thermocouple is attached at the back face using high temperature cement [19]. Location of the required heat  \\x0c', \"272  A . Nisar et al.  / Carbon 111 (2017) 269e282  Fig. 4. Bright field TEM micrographs of the ZrB2-based composites (a) pure ZrB2 (b) Z20S (c) Z10C (inset shows the buckling and bending of CNT after SPS processing) and (d) Z20S10C showing SiC and CNT sits at the inter-granular junctions. Selected area diffraction (SAD) pattern showing (e) hexagonal ZrB2 (f) hexagonal SiC and (g) nanocrystalline CNT. (A colour version of this figure can be viewed online.)  plasma arc jet test was calculated as an average of area fraction of (001), (100) and (101) XRD peaks. The crystallite size (t) of ZrB2based composites before and after plasma arc-jet test was evaluated using Scherrer's equation:  V-I 51. Diametrical compression test was carried out using 100 kN Universal Testing Machine (UTM, Model: BiSS Ltd.) at a strain rate of 0.05 mm/min for all the SPS processed composites. The fracture strength (sf) of the material is calculated using formula [22,23]:  t ¼ 0:9 l b cosq  q ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi  s  sf ¼ 2 P phd  (4)  (5)  m \\x00 b2 b2  where, b ¼ , bm and bs are full-width half maximum (FWHM) of the most intense peak of the sample and that of the standard Al2O3 sample respectively, l is the wavelength of incident X-ray, and q is Bragg's diffraction angle. Further, to study the damage of CNT after the plasma arc jet exposure, micro-Raman spectroscopy (Princeton Instruments, STR Raman, TE-PMT detector) was carried out using Nd-YAG green laser (l ¼ 532 nm) with laser power of 12.5 mW, an objective of 50\\x02 and spatial resolution \\x001 in the backscattered mode. The surface morphology of of 0.5 cm ZrB2-SiC-CNT composites (SPS and after plasma exposure) and profile of oxide layer after plasma exposure along the cross section were imaged using W-SEM (JEOL JSM-6010LA and Zeiss Model EVO 50). Transmission electron microscopy (TEM, FEI UT 20) of the SPS processed ZrB2-SiC-CNT composites and the oxidized Z20S10C sample (after scratching the surface oxide) was performed at an operating voltage of 200 kV. Hardness and elastic modulus of the as-sintered compacts and on the cross-section of the oxide layer were measured by using instrumented micro indenter (MHT, CSM instruments, Switzerland) on the polished samples. The load function comprises loading to 2 N at rate of 4 N/min and holding at maximum load for 10 s followed by unloading at the same rate using Vickers' indenter of type  where, P is the maximum load at which fracture occurs, h is the thickness and d is the diameter of the sample. Three samples were tested and an average value (obtained within 15% deviation) is reported.  3. Results and discussion  3.1. Densification of ZrB2-SiC-CNT composites  The densification of ZrB2-based composites increased from 93% for pure ZrB2 to ~100% with synergistic reinforcement of both SiC and CNT in ZrB2 under similar SPS processing conditions, presented in Fig. 2 (also see Table 1). During SPS, the recorded punch displacement profile reveals three distinct regions, i.e. initial compaction (step I) due to the rearrangement of powder particles; intermediate step (step II) corresponding to thermal expansion, and shrinkage (step III) corresponding to final densification as shown in Fig. 2. Since, step III corresponds to shrinkage, is a crucial step for the final densification, which is observed to increase with SiC and CNT reinforcement [20].  \\x0c\", \"A . Nisar et al.  / Carbon 111 (2017) 269e282  273  contraction of the ceramic matrix during sintering [6,20,24]. The ratio of the intensity of Dto G-peak (ID/IG) increased from 0.97 for pristine CNT powder to 0.99 (for Z10C) and 1.03 (Z20S10C) indicating damage of CNT structure (bending and buckling in CNT, shown later in Fig. 4).  3.3.  Structural stability of CNT during SPS processing  Structural stability of CNT is essential in rendering enhanced toughness, so it is imperative to investigate the interfaces and structure of various phases of the SPS composite pellets (see Figs. 4aed). The absence of impurities (oxide, glassy phase etc.) at triple junctions and presence of dislocation tangles even in pure ZrB2 (Fig. 4a) suggests removal of oxide impurities during SPS, which enables solid-state sintering. It can be well accentuated from Figs. 4b and c that SiC and CNT are tightly encapsulated within ZrB2 matrix. However, the bending and buckling of CNT can also be seen in the inset of Fig. 4c, which microstructurally supports the increment in ID/IG ratio obtained during Raman analysis (Fig. 3b). Again, along with the presence of anisotropic ZrB2 grains, Fig. 4d shows the presence of both SiC and CNT, which is further confirmed by the selected area diffraction (SAD) patterns as shown in Figs. 4eeg, respectively. The SAD pattern for hexagonal ZrB2 (Fig. 4e) and 6HSiC (Fig. 4f) indexed corresponding to the zone axis [111]. The nanocrystalline nature of the CNT is evident in Fig. 4g.  3.4. Phase and microstructural characterization after plasma arcjet testing  Preand post-oxidation investigation of ZrB2-SiC-CNT composites showed no erosion on the exposed face, which elicits that all composites survived the flux of 2.5 MW/m2 for 30 s, see Fig. 5. The test (with flux of 2.5 MW/m2) was enough to simulate thermal shock, which apparently is withstood by all the samples as the oxide scale was adherent (no occurrence of spallation or microcracks) for all the composites (see Fig. 5a). Apparently, surfaces of ZrB2 and Z10C have turned white due to the oxidation at high temperature, however, no considerable changes were visible in Z20S and Z20S10C. The back face rise in temperature for all plasma arc-jet tested UHTC composites (shown in Fig. 5b) is observed to decrease with SiC addition and increases with CNT reinforcement in the ZrB2 matrix. Oxidation of ZrB2-SiC-CNT composites during plasma arc jet exposure has been confirmed from the subsequent appearance of monoclinic ZrO2 and SiO2 peaks obtained in XRD pattern (Fig. 6a), however, the retention of ZrB2 phase was observed and further quantified (on the surface oxide) as an average of area fraction of ZrB2 under XRD peaks, see Table 2. It is elucidated that the formation of surface oxide was suppressed with SiC and CNT addition as the amount of retained ZrB2 phase was increased from ~30% (for ZrB2) to 47% (for Z20S), 38% (for Z10C) and 42% (Z20S10C). The crystallite size (t) evaluated using Scherrer's formula (Eq. (4)) for all as-prepared samples are of similar order (~39e44 nm), however, the crystallite size noticeably increased after plasma arc-jet exposure (~53e62 nm), shown in Table 2. Raman spectrum showed two distinctive peaks at 1348.3 and 1573.6 cm for Z10C and 1349.2 and 1588.6 cm for Z20S10C (Fig. 6b), which corresponds to Dand G-peaks of carbon, respectively [20]. Thus, the retention of CNT is confirmed in SPS processed ZrB2-SiC-CNT composites even after plasma arc jet exposure. Crosssectional Raman line analysis across the oxide scale shown in Fig. 6c (along a line with distinct point, each at a distance of 5 mm) of Z20S10C samples is presented in Fig. 6d. It is elicited from the graph (Fig. 6d) that ZrB2, being Raman in-active, showed no peaks.  \\x001  \\x001  Fig. 5. (a) Photographs of the SPS pellets before and after plasma-arc jet testing and (b) Back face temperature profiles of ZrB2-SiC-CNT composites with time when subjected to plasma arc jet for 30 s. (A colour version of this figure can be viewed online.)  3.2. Phase retention in spark plasma sintering of ZrB2-SiC-CNT composites  X-ray diffraction (XRD) patterns for powder mixtures as well as SPS sintered ZrB2-based composites are presented in Fig. 3a. The presence of characteristic peaks of ZrB2 and SiC elicit retention of phases without any undesirable interfacial reactions after SPS processing. However, the intensity of SiC phase was low due to poor X-ray reflecting the ability of SiC [20]. No peaks corresponding to CNT were detected in XRD, whereas, Raman spectra (Fig. 3b) elucidate Dand G-peaks confirming the retention of CNT even after SPS processing. The shift in the Dand G-peaks in the SPS samples when compared to pristine CNT (D-peak at 1351.8 cm \\x001) elicits that the compressive stress was and G-peak at 1584.4 cm introduced after SPS processing, possibly due to thermal  \\x001  \\x0c\", '274  A . Nisar et al.  / Carbon 111 (2017) 269e282  Fig. 6. (a): XRD spectra of plasma arc jet exposed ZrB2-SiC-CNT composites pellet (b) Top surface Raman spectra of Z20S and Z20S10C samples, (ced) Optical image and corresponding cross-sectional Raman spectra of Z20S10C sample showing the presence of graphitic peaks even after plasma arc jet exposure. Orange coloured line indicates the direction in which Raman line analysis has been carried out. (A colour version of this figure can be viewed online.)  Table 2 Crystallite size, retention of ZrB2 phase, the thickness of the oxide layer, and oxidation rate after plasma arc-jet exposure of ZrB2-SiC-CNT composites.  Sample ID  Crystallite size (t) of ZrB2 in SPS pellet (nm)  Retention of ZrB2 phase after exposure  Thickness of the oxide layer (mm)  Linear oxidation rate (mm/s)  Before plasma arc jet  After plasma arc jet  ZrB2 Z20S Z10C Z20S10C  44 ± 2 39 ± 1 40 ± 2 40 ± 1  57 ± 2 62 ± 2 53 ± 2 55 ± 1  ~30% ~47% ~38% ~42%  22.9 ± 0.3 10.1 ± 0.1 16.8 ± 0.3 13.2 ± 0.3  0.77 0.34 0.56 0.44  However, Dand G-peaks were observed wherever CNT was present both in the oxide scale as well as the unaffected sample. \\x001, Correspondingly, the presence of oxide bands at 630.9 cm \\x001 and 1011.4 cm \\x001 is mainly due to a vibrational mode 959.1 cm involving Si-O-Zr linkages were also confirmed in the samples after  plasma arc jet exposure [25]. Among these, the Raman peak at \\x001 corresponds to the break-up of Si-O-Zr linkages due to 959.1 cm the exsolution of ZrO2 as tetragonal zirconia [25]. Of the various reaction products mentioned in the reaction schemes (Eqs. (6)e(10)), ZrO2 is stable at > 2000 \\x0e C with a melting  \\x0c', 'A . Nisar et al.  / Carbon 111 (2017) 269e282  275  small pits throughout the sample. The cross-sectional SEM micrographs of plasma exposed samples elicit the thickness (presented in Table 2), morphology and composition (see appendix 1) of oxide scale (ZrO2), shown in Figs. 9aed and Figs. 9eeh respectively, however, the composition of oxide scale is presented in appendix 1. It is to be mentioned here that the oxide scale spalled off from ZrB2 (Fig. 9a) during the preparation of cross-sectional sample (sectioning followed by hot mounting and ultra-sonication). The thickness of oxide scale decreased with SiC and CNT reinforcement from 23 mm (in ZrB2, Fig. 9a) to 10 mm (in Z20S, Fig. 9b), 17 mm (in Z10C, Fig. 9c) and 13 mm (in Z20S10C, Fig. 9d). It is evident from the microstructure (Figs. 9eeh) that the temperature experienced by the composites is in excess of 2715 \\x0e C, i.e. the melting temperature of ZrO2 [26]. The linear oxidation rate (calculated using Eq. (1)), presented in Table 2 elicits that the oxidation rate decreases from 0.77 mm/s for pure ZrB2 to 0.34 mm/s (for Z20S), 0.56 mm/s (for Z10C) and 0.44 mm/s (for Z20S10C) has been attributed to the high thermal conductivity (discussed in the following section) with SiC and CNT reinforcement. TEM of plasma arc jet exposed Z20S10C sample showed the presence of CNT and transformed CNT (partially graphitized) as shown in Fig. 10a, which supplement the results obtained in Raman spectroscopic analysis (Figs. 6bed) and microstructure obtained before plasma exposure (Figs. 8c and d). The morphology of the oxide scales of SiO2 and ZrO2 is elucidated from Figs. 10b and c. The SAD pattern in Figs. 10def corresponds to the zone axis [111] for monoclinic ZrO2, triclinic SiO2, and CNT, respectively. The survival of CNT (along with damaged CNT), even after plasma arc jet exposure, insinuates the plausibility of CNT in resisting oxidation of ZrB2 by both physical (Fig. 10 a) and chemical means (Eqs. (8) and (10)) as explained in the following section.  3.5. Oxidation behavior  3.5.1.  Thermal conductivity  The thermal conductivity of ZrB2-SiC-CNT based UHTC was determined based on the values calculated from the correlation function of thermal diffusivity, specific heat capacity and density (Eqs. (2) and (3)) are presented in Table 3. Thermal conductivity for pure ZrB2 (~48.9 W/mK) is well in accordance with that reported in the literature [8], and is observed to increase with synergistic reinforcement of SiC and CNT (Z20S10C) to 61.8 W/mK. As the temperature increases (from 50 to 1200 \\x0e C), defect concentration in the material increases, which leads to phonon scattering (known as Umklapp scattering) and, thus, a decrease in the thermal conductivity is observed. The addition of SiC and CNT increases the thermal conductivity based on the dispersed phase models [14,19], providing a conductance path (along the grain boundary, see Fig. 6) for heat dissipation, thereby, providing improved oxidation resistance (the high bow shock associated with plasma arc-jet is enough to generate thermal shock). However, a little decrease in thermal conductivity values, in case of CNT reinforced composites (Z20S10C, when compared with Z20S), is attributed to the distribution and alignment of CNT because CNT reinforcement shows higher thermal conductivity when it is aligned in one direction) [18]. The measured thermal conductivity values are higher than for some alloys and ceramic matrix composites (CMCs, for example: at 300 K, l ~ 13.6 W/mK for Inconel 617 and l ~ 14.5 W/mK for a standard C/ C-SiC [30]. The rise in back face temperature (see Table 3, and Fig. 5b) is attributed to the thermal conductivity, porosity content, and emissivity of the material. Rise in temperature (Tb) of ZrB2 (1034 \\x0eC) in spite of lower thermal conductivity as compared to reinforced composites is due to high porosity content (~7%) in the material as  Fig. 7. Plot of standard free energy change vs absolute temperature using FactSage [28] for the reactant and product. (A colour version of this figure can be viewed online.)  point of ~2715 \\x0eC [26], also revealed from the XRD results. B2O3 has a low melting point (450 \\x0e C) and high vapor pressure and hence it vaporizes at temperatures above 1100 as per the below mentioned reactions [27].  \\x0e C  2  5  ZrB2 ðsÞ þ O2 ðgÞ/2  5  ZrO2 ðsÞ þ 2  5  B2O3 ðlÞ  B2O3 ðlÞ/B2O3 ðgÞ  ZrO2 ðsÞ þ B2O3 ðlÞ þ 5C ðsÞ/ZrB2 ðsÞ þ 5CO ðgÞ  2  3  SiC ðsÞ þ O2 ðgÞ/2  3  SiO2 ðsÞ þ 2  3  CO ðgÞ  SiO2 ðsÞ þ 3C ðsÞ/SiC ðsÞ þ 2CO ðgÞ  (6)  (7)  (8)  (9)  (10)  The plot of standard free energy change against absolute temperature (using FactSage [28]) for reactants and products in the aforementioned reactions (see Fig. 7), from which it is inferred that temperatures \\x15 1196.5 \\x0e C, 575.8 \\x0eC, Eqs. (6)e(10) are feasible at 1499.4 \\x0e C, 1222.6 \\x0eC and 1452.4 \\x0eC respectively under the pressure \\x0015 atm (during re-entry in real-life application [29]). It is of pO2 10 anticipated that the oxidation of ZrB2 to ZrO2 (Eq. (6), exothermic) and SiC to SiO2 (Eq. (9), exothermic) are responsible for the net weight gain (later explained in section 3.5.2) while Eq. (7) (endothermic) result as the weight loss. However, from Eqs. (8) and (10) (endothermic), it is elicited that the addition of CNT suppresses the formation of oxide scales (i.e. ZrO2 and SiO2), which is also evident from the XRD analysis (Table 2). Typical SEM images depicting composites before and after oxidation, are shown in Fig. 8. Figure 8a shows the surface morphology of as processed ZrB2 with porosity (~7%) while the presence of SiC and CNT (darker spots, Figs. 8bed) appears to be uniformly distributed in ZrB2 matrix. Formation of multiple cracks and deep pits seen on ZrB2 sample (Fig. 8e) are attributed to the escape of B2O3 (g) phase (see Eq. (7)), resulting in pore formation which accelerates oxidation diffusion. In contrast, a few traces of ZrO2 with the absence of cracks was observed in Z20S (Fig. 8f). Z10C and Z20S10C samples in Figs. 8g and h showed fewer cracks and  \\x0c', '276  A . Nisar et al.  / Carbon 111 (2017) 269e282  Fig. 8. SEM micrographs of SPS processed ZrB2-based composites (a) ZrB2, (b) Z20S, (c) Z10C and (d) Z20S10C and (eeh) display cracks, pits and oxide scale obtained after plasma arc jet exposure. (A colour version of this figure can be viewed online.)  per the effect of phonon-pore scattering on thermal conductivity [18]. Conversely, the maximum rise in temperature in Z10C (1081 \\x0e C) is due to the high thermal conductivity of the CNT, however, composites with SiC (i.e., Z20S (853.2 \\x0eC) and Z20S10C (842.3 \\x0eC)) showed the least rise in temperature due to their higher surface emissivity (~0.98) [31]. Also, composites with SiC reinforcement forms oxide scale of SiO2-ZrO2 on the surface (passive oxidation), which hinders the rise in back face temperature (Tb) due to the poor thermal conductivity of oxide.  3.5.2.  Thermal studies using TGA/DSC  Thermal studies using TGA (although, conditions are not as severe as that of plasma arc jet exposure) have been carried out to provide an insight to the oxidation mechanisms occurring in ZrB2SiC-CNT composites. The final increase in the weight of the composites are: 24%, 5%, 18% and 2%, respectively, for ZrB2, Z20S, Z10C and Z20S10C (see Table 4) illustrate that the resistance to oxidation increases with SiC and CNT addition (see Fig. 11 a) due to the formation of SiO2 (Eq. (9)) and sealing mechanism offered by CNT [19]. Figure 11a elicits the shift to higher onset temperature of oxide  \\x0c', 'A . Nisar et al.  / Carbon 111 (2017) 269e282  277  Fig. 9. Back-scattered SEM micrographs showing the thickness of the oxide scale of (a) ZrB2, (b) Z20S, (c) Z10C and (d) Z20S10C and (eeh) are details of the oxide scale (secondary electron mode) in the respective composite. Double-headed arrow (in Figs. aed) shows the oxide scale thickness, while single-headed arrows (in Figs. eeh) represent the melting and re-solidified grains. (A colour version of this figure can be viewed online.)  formation after which a rapid weight change is observed (Table 4), i.e. 718.5 \\x0eC, 697.1 \\x0eC and 706.3 \\x0eC for Z20S, Z10C, and Z20S10C, respectively, when compared to that of pure ZrB2 (679.3 \\x0e C). The oxidation products (B2O3 and SiO2, in liquid form) formed during the initial stage can penetrate into the pellet through surface pores [32] and offers good ablation resistance [32,33] when compared to  that of pure ZrB2. Figure 11b showed an exothermic peak at ~652.2 \\x0eC for pure ZrB2, which shift towards higher temperature i.e., 661.9 \\x0eC, 676.1 \\x0eC and 690.5 \\x0eC respectively for Z20S, Z10C, and Z20S10C with SiC and CNT addition. The presence of an endothermic peak at around 1350 \\x0e C, observed in case of ZrB2 and Z10C suggest a possible  \\x0c', '278  A . Nisar et al.  / Carbon 111 (2017) 269e282  Fig. 10. TEM micrographs of the plasma exposed Z20S10C sample showing (a) oxide scale and the presence of CNT (represented by single-headed arrows), (b) SiO2 particle (inside feature represents twin boundary), (c) ZrO2 particle and (def) SAD pattern confirming the crystal structure of the ZrO2, SiO2 and retained CNT after plasma arc jet exposure. (A colour version of this figure can be viewed online.)  Table 3 Thermal diffusivity, specific heat capacity, thermal conductivity and back face rise in temperature of ZrB2-SiC-CNT based composites.  Sample  Thermal properties of ZrB2-based composites at different temperatures 50 \\x0e C 300 \\x0e C 600 \\x0e C  ZrB2 Z20S Z10C Z20S10C  D  0.236 0.276 0.204 0.240  Cp  340.4 446.0 415.0 485.2  \\x001, Cp in J kg Where, D in mm2 s  k  D  Cp  k  D  48.9 0.180 415.9 45.4 0.147 67.6 0.197 623.5 67.6 0.173 48.9 0.159 549.7 49.8 0.133 61.8 0.177 624.3 58.7 0.160 \\x001, k in W/mK and Tb ¼ back-face temperature.  \\x001 K  Cp  512.7 668.5 744.8 799.0  k  45.7 65.2 56.2 67.4  900 \\x0e C  D  0.131 0.151 0.130 0.127  Cp  536.7 732.8 632.6 802.4  k  42.7 60.5 46.6 51.4  1200 \\x0e C  D  0.125 0.121 0.130 0.121  Cp  559.7 761.2 671.3 816.7  k  42.3 50.6 49.5 52.3  Tb (\\x0e C)  1034.0 853.2 1081.3 842.3  Table 4 Thermal studies of ZrB2-SiC-CNT composites up to 1500 \\x0e C in oxygen atmosphere.  Composition  % wt. gain  Onset rapid weight change temperature (\\x0e C)  Peak temperature of oxide formation (\\x0e C)  ZrB2 Z20S Z10C Z20S10C  24 5 18 2  679.3 ± 6.4 718.5 ± 3.8 697.1 ± 3.7 706.3 ± 2.1  652.2 ± 2.8 661.9 ± 4.4 676.1 ± 2.9 690.5 ± 3.1  Enthalpy of oxide formation DH (kJ/g) \\x001.6 ± 0.1 \\x000.3 ± 0.1 \\x000.8 ± 0.2 \\x000.6 ± 0.1  transformation of m / t ZrO2 [10]. Based on this, the efficacy of SiC reinforcement in protecting ZrB2 from oxidation and the absence of transformation of m / t ZrO2 is well elucidated. The enthalpy (DH) of formation of oxide decreases from 1.6 kJ/g for pure ZrB2 to 0.3 kJ/ g, 0.8 kJ/g and 0.6 kJ/g for Z20S, Z10C and Z20S10C, respectively, which insinuates an increase in the thermal stability with SiC and CNT reinforcement.  3.5.3.  Effect of SiC and CNT reinforcement on oxidation mechanism  The schematic illustration of the factors that govern the oxygen transport mechanism (SiC and CNT reinforcement; processing induced defects) is shown in Figs. 12aec. The presence of dislocation network in ZrB2 grain (see TEM image, Fig. 12d) act as an active site, which triggers the diffusion of oxygen via grain and grain boundaries (GBs) [34]. In the case of SiC reinforced composites,  \\x0c', 'A . Nisar et al.  / Carbon 111 (2017) 269e282  279  composites, the mechanical properties of the as-processed and after plasma arc-jet tested ZrB2-SiC-CNT composites were evaluated using instrumented indentation technique. The hardness and Young\\'s modulus, tabulated in Table 5 are calculated from the load vs displacement graph (shown in Fig. 13) using Oliver-Pharr indentation method for ZrB2-based composites before and after plasma arc jet exposure. Vickers\\' hardness for as-processed ZrB2 is increased from 13.2 GPa to 19.3 GPa for Z20S, 18.5 GPa for Z10C and 21.0 GPa for Z20S10C, attributed to the higher relative densities of composites after SPS processing. It may be noted that plasma arc jet exposed samples showed lower hardness than that of unexposed samples. However, a similar trend of hardness was maintained even after plasma exposure where it increases from 6.2 GPa for pure ZrB2 to 12.0 GPa (for Z20S), 10.2 GPa (for Z10C) and 12.8 GPa (for Z20S10C). Z20S10C showed the highest hardness values elucidating the synergy of SiC and CNT reinforcement in ZrB2 matrix before and after plasma arc jet exposure. The measured Young\\'s modulus for the as-processed ZrB2 (~93% dense) ~378 GPa is in agreement with the reported value [36]. This value increases to 405 GPa for Z20S, 401 GPa for Z10C and 411 GPa for Z20S10C when compared with pure ZrB2. Correspondingly, the modulus decreases to 212 GPa for monolithic ZrB2 (after plasma arc jet exposure) which then increases with SiC and CNT reinforcement to 331 GPa (for Z20S), 288 GPa (for Z10C) and 342 GPa (for Z20S10C) after plasma arc jet exposure. Based on the hardness and modulus of these composites, the H3/E2 ratio was determined (see Table 5), which confirm a higher resistance to plastic deformation with SiC and CNT reinforcement in ZrB2 before and after plasma arc jet exposure when compared to that of ZrB2. The load versus displacement relationship during the diametrical compression test (see appendix 2) is used to calculate the fracture strength of as processed ZrB2-based composites (see Table 5). The fracture strength obtained for pure ZrB2 (32.6 MPa) is in agreement with the reported value [8]. The fracture strength increased to 78 MPa (2.5 times for Z20S), 54.3 MPa (1.8 times for Z10C) and 81.1 MPa (2.7 times for Z20S10C) when compared to pure ZrB2. The increased strength with reinforcement indicates higher load bearing capacity of composites via CNT bridging and interfacial bonding of matrix with reinforcement [20].  3.5.5. Effect of residual stress on structural stability of ZrB2-SiC-CNT composites  There are several strengthening mechanisms that occur in composite materials: (i) crack bridging (ii) micro-cracking when the misfit between coefficient of thermal expansion (CTE) introduces a stress field, and (iii) residual stress toughening due to difference in CTE of the matrix and reinforced particles, which creates a local compressive stress field in the matrix, thus, decreasing the stress intensity factor. Herein, authors have isolated the contributions of thermal residual stresses in composites, which arise due to the coupling of different phases with different thermoelastic properties. Hence, the strength of the ZrB2 is governed by the reinforcement of SiC and CNT (ignoring the effect of distribution and size) and the generation of interfacial residual stresses (computed analytically) in the material. It may be mentioned that the extrinsic effects (such as glassy phase content, micro-cracking, porosity, etc.) have not been accounted for in estimating residual stresses at the interface. The linear CTE of a two-phase composite (assuming that each phase is isotropic), computed by Levin [37], Rosen and Hashin [38], is provided as:  \"  #  \\x10  \\x11  aeff  u  ¼ fmam þ fr ar þ ðam \\x00 ar Þ  \\x00 1  Kr  1  Km  1  Keff  \\x00 1  Km  \\x00 1  Kr  (13)  Considering that Kl \\x14 Keff \\x14 Ku, the upper and lower bounds of K  Fig. 11. Thermal analysis at 1500 \\x0e C on ZrB2-SiC-CNT based composites: (a) TGA and (b) DSC. (A colour version of this figure can be viewed online.)  oxygen can transport via ZrB2, SiC particles, and GBs (see Fig. 12b), however, transport through GBs is favored over particles [19]. Also, SiC is usually observed above 1700 \\x0eC, the rapid oxidation of therefore, it protects ZrB2 by the formation of passive SiO2 (see Fig. 12b), further explored by the negligible oxidation observed till 1500 \\x0e C (see Fig. 11a). Moreover, during thermal studies the presence of low angle grain boundaries (LAGBs) at the interface of ZrB2 and SiC (TEM image, Fig. 12e) act as dormant site (having low energy) and, thus, provides oxidation resistance [35]. Similarly, the survival of CNT from harsh plasma jet exposure (shown earlier in Figs. 6 and 10) establish restriction of oxidation of ZrB2 via grain sealing mechanism (Figs. 12c and f) [19]. Oxygen diffusion through interfaces may be preferred, thus, CNTs, sealing the grains, may be sacrificially oxidized and protect the underlying ZrB2 and SiC (see Eqs. (8) and (10)). Succinctly, the synergy of SiC and CNT reinforcement ascertain ZrB2 as a potential structural material with high damage tolerance.  3.5.4. Evaluating mechanical performance of ZrB2-SiC-CNT composites  To further  investigate  the  structural  integrity of ZrB2-based  \\x0c', '280  A . Nisar et al.  / Carbon 111 (2017) 269e282  Fig. 12. Schematic illustration of oxygen transport in (a) ZrB2 matrix, (b) with SiC reinforcement (c) grain sealing hindering diffusion of oxygen with CNT addition and (def) corresponding TEM image eliciting the microstructural features (after SPS processing) affecting oxidation kinetics. Here, GB ¼ grain boundary (in Figs. aec) and LAGBs ¼ low angle grain boundaries (in Fig. e). (A colour version of this figure can be viewed online.)  Table 5 Comparison of mechanical properties of as-prepared and plasma arc jet exposed ZrB2-SiC-CNT composites using instrumented indentation technique.  Sample ID  Before oxidation  ZrB2 Z20S Z10C Z20S10C  Hv (GPa)  13.2 ± 1.0 19.3 ± 0.6 18.6 ± 0.8 21.0 ± 0.7  E (GPa)  378 ± 5 405 ± 6 401 ± 10 411 ± 7  H3/E2 (MPa)  sf (MPa)  16.1 43.8 40.1 54.8  32.6 78.0 54.3 81.1  After oxidation  Hv (GPa)  6.2 ± 1.4 12.0 ± 0.6 10.2 ± 0.9 12.8 ± 0.5  E (GPa)  232 ± 11 331 ± 9 288 ± 8 342 ± 2  H3/E2 (MPa)  3.1 15.8 12.8 17.9  can be obtained from Hashin-Shtrikman bounds [39]. Under such conditions, the above Eq. (13) is further modified to calculate both the upper and lower bounds of a in the following manner:  \\x10  \\x11  aeff  u  \\x10  aeff  \\x11  l  ¼ am \\x00 fr ðam \\x00 ar Þ  Kr ð3Km þ 4Gm Þ Km ð3Kr þ 4Gm Þ þ 4fr Gm ðKr \\x00 km Þ  ¼ ar þ fm ðar \\x00 am Þ  Km ð3Kr þ 4Gr Þ Kr ð3Km þ 4Gr Þ þ 4fmGr ðKm \\x00 kr Þ  (14)  (15)  where, ðaeff Þu and ðaeff Þl are the upper and lower bounds of CTE  a  \\x006  respectively of a given composite whereas a, K, G and f are CTE, bulk modulus, shear modulus and volume fraction with subscript “m” and “r” set for matrix and reinforcement respectively. The values of n (for ZrB2 ¼ 0.17, SiC ¼ 0.14 and CNT ¼ 0.17, calculated using ROM), ZrB2 ¼ 6.8 \\x02 10 \\x001, SiC ¼ 3.5 \\x02 10 (for K K and CNT ¼ 2.5 \\x02 10 \\x006 K \\x001), K (for ZrB2 ¼ 229 GPa, SiC ¼ 234 GPa and CNT ¼ 190 GPa) and G (for ZrB2 ¼ 211 GPa, SiC ¼ 41 GPa and CNT ¼ 150 GPa) used for calculations (are taken from the literature [19,20,40]). The quantification of such thermal residual stresses in the composites becomes primarily important, since the processing has been carried out at a very high temperature where these stresses develop (DT). Since the composites are being cooled from the final densification temperature (1650 \\x0e C) to room temperature; ZrB2  \\x006  \\x001  \\x0c', \"A . Nisar et al.  / Carbon 111 (2017) 269e282  a ¼  0  Em  ð1 \\x00 nm Þε  281  (20)  The stability of the oxide scale formed also depends on internal stresses developed during plasma arc jet exposure. After the plasma arc jet exposure, both the oxide and the remaining composite contract upon cooling [37], leading to compressive strain in oxide scale and tensile strain in the remaining composite. So, the oxide scale relaxes (as the cooling is very rapid) by generating cracks on the surface as observed in Fig. 8. It is elucidated from Table 6 that the addition of secondary phase (SiC and CNT) introduce an interfacial stress (compressive in the reinforcement and tensile in the matrix), which provides enhanced structural integrity and toughness [20]. The values of compressive stress ranges from 34.3 MPa (for Z20S) to 43.1 MPa (for Z10C) and 39.4 MPa (for Z20S10C), but are only ~5% of that tensile stress in the matrix. It is observed that the residual tensile stress in the matrix decreases with reinforcement of SiC and CNT (Table 6, Taya's model). The biaxial model follows the same trend of residual stress with SiC and CNT reinforcement. Thus, CNT-reinforced composites (Z10C and Z20S10C) showed structural stability (due to a reduction in residual tensile-stresses) along with the thermal stability established by thermal studies (discussed earlier in section 3.5.2.). Hence, enhancement in the interfacial residual compressive stresses with the synergistic addition of SiC and CNT provide an assessment of the structural integrity, which is mandatory for aerospace applications.  4.  Conclusions  ZrB2-based UHTC composites were spark plasma sintered achieving enhanced densification from 93% to ~99.8% with synergistic reinforcement of SiC and CNT. Oxidation behavior of ZrB2-SiCCNT composites was evaluated through plasma arc jet with 30 s exposure under heat flux of 2.5 MW/m2. The decrease in the linear oxidation rate from 0.77 mm/s (for ZrB2) to 0.34 mm/s (for Z20S), 0.56 mm/s (for Z10C) and 0.44 mm/s (for Z20S10C) is attributed to the enhanced thermal conductivity with SiC and CNT addition. The structural integrity of these UHTC composites was measured in terms of hardness which increases up to 1.6 times (13.2e21.0 GPa) before and up to 2.1 times (6.2e12.8 GPa) with synergistic reinforcement of SiC and CNT in ZrB2 after exposure to harsh aerothermal conditions. Raman spectroscopic analysis revealed that CNT survived the extreme oxidising environment, further supplemented by TEM analysis of Z20S10C. In addition, enhanced thermal conductivity with SiC and CNT addition in ZrB2 indicates an improved oxidation resistance via heat dissipation mechanism. The the oxide formation (from 679 \\x0eC to high temperature shift of 706 \\x0eC) and decrease in the enthalpy of oxide formation (from 1.6 to 0.6 kJ/g), elicits enhanced thermal stability and oxidation resistance with synergistic reinforcement of SiC and CNT in ZrB2. The synergy of SiC and CNT in providing oxidation resistance (protective SiO2  Fig. 13. Load-displacement curves for ZrB2-SiC-CNT composite before and after plasma arc jet exposure. (A colour version of this figure can be viewed online.)  (6.8 \\x02 10 \\x006 K \\x001) tends to shrink faster than the reinforced particles (SiC, 3.5 \\x02 10 \\x006 K \\x001 and CNT, 2.5 \\x02 10 \\x006 K \\x001) leading to tensile stress state in the matrix, sm, and corresponding compressive stress in the reinforced particles, sr, as evaluated using Taya's model [41], and tabulated in Table 6.  sr ¼ \\x00ð1 \\x00 f Þε  0  f  sm  sm ¼ þEm  0  2f bε  ð1 \\x00 f Þðb þ 2Þð1 þ nm Þ þ 3bf ð1 \\x00 nm Þ  b ¼ 1 þ nm 1 \\x00 2nr  Er Em  0 ¼ ðar \\x00 am ÞDT  ε  (16)  (17)  (18)  (19)  where, Er, Em, nr and nm are the Young's modulus and Poisson's ratio (for ZrB2 ¼ 0.17, SiC ¼ 0.14 and CNT ¼ 0.17, calculated using ROM [19,20,40]) of the reinforcement and matrix respectively; ε is the thermal expansion misfit strain and DT is the temperature at which as 1400 \\x0e C) stresses begin to accumulate (set [42,43]. Young's modulus for the reinforced phase, is estimated by ROM, considering modulus of matrix (ZrB2) as Em ¼ 378 GPa and modulus Ec for each composite is both from values presented in Table 4. Correspondingly, the developed biaxial residual stress (s) obtained from the relation:  Er  is  0  Table 6 Theoretical calculation of the coefficient of thermal expansion (CTE), modulus and residual stresses.  Sample ID  CTE (\\x0210  \\x006/K)  Rule of mixture (ROM)  Hashin-Shtrikman model  ZrB2 N.A. Z20S 6.14 Z10C 6.37 Z20S10C 5.71 Where, N.A ¼ not applicable.  Upper bound  Lower bound  N.A. 6.15 6.35 5.69  N.A. 6.13 6.30 5.62  Modulus (GPa)  Residual stress (MPa)  Ec  378 405 401 411  Er  N.A. 513 608 488  Taya's model  Biaxial  sm  N.A. 630.1 352.4 428.0  sr  N.A. \\x0034.3 \\x0043.1 \\x0039.4  N.A. 520.1 347.6 481.7  \\x0c\", '282  A . Nisar et al.  / Carbon 111 (2017) 269e282  formation by SiC and grain sealing by CNT) and enhanced thermal \\x0e C) conductivity (from 42.3 W/mK to 52.3 W/mK at 1200 is corroborated. These findings demonstrate that the interplay of reinforcement affects the generation of residual stresses at the interface (analytically quantified), which, then, governs the strength of ZrB2-SiC-CNT composites, making them a suitable candidate for application in high-damage tolerant structures.  Acknowledgements  Authors at Indian Institute of Technology Kanpur (IITK) acknowledge the financial support received from IITK-Space Technology Cell (Vikram Sarabhai Space Center, Indian Space Research Organization (ISRO), Trivandrum, India). Mr. Vincent Xavier (Technical Staff) and Mr. Fazil Mohammad are acknowledged for helping with plasma-arc jet exposure testing and thermal conductivity measurements, respectively, at ISRO. KB acknowledges P.K. Kelkar fellowship, IITK. Authors also acknowledge TEM Facility, MSE Department, IITK and the support from the Advanced Centre for Materials Science (ACMS) at IITK for extending characterization facilities (SEM, mechanical testing and instrumented indentation testing).  Appendix A.  Supplementary data  Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.carbon.2016.10.002.  References  [4]  for  vol. X,  Ceramists,  lightself [1] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, High-strength zirconium diboride-based ceramics, J. Am. Ceram. Soc. 87 (6) (2004) 1170e1172. [2] A.E.M. (Ed.), Data Collected from Phase Diagrams American Ceramic Society, Westerville, OH, 1994. [3] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (5) (2007) 1347e1364. B. Xu, S. Zhou, C. Hong, J. Han, X. 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},{
  "_id": 42,
  "PDF": "Effect of HfB2 and WC additives on the ablation resistance of ZrB2–SiC composite coating manufactured by SPS.pdf",
  "Text": "['Ceramics International 46 (2020) 25106-25112  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Effect of HfB2 and WC additives on the ablation resistance of ZrB2-SiC composite coating manufactured by SPS  T  S. Ali Akbarpour Shalmania, M. Sobhania,∗, O. Mirzaeea, M. Zakerib,∗∗  a Faculty of Materials and Metallurgical Engineering, Semnan University, Semnan, b Ceramic Dept, Materials and Energy Research Center, Karaj, Iran  Iran  A R T I C L E  I N F O  A B S T R A C T  ZrB2-SiC composite coating with HfB2-WC as additives was successfully manufactured on graphite substrate by spark plasma sintering method. The effect of HfB2 and WC on coating properties was investigated. The microstructure, phase changes and ablation resistance of the coatings were studied. Results showed that composite coatings with the thickness of 500 μm and penetration depth of 250 μm were obtained for all samples. Results of the oxyacetylene flame test showed that the ablation resistance of the coated graphite was significantly increased. The minimum ablation weight percentage and rate of 2% and 3 mg s−1 were obtained for the composite coating with 3.75% of each additive.  Keywords: Ablation Composite SPS ZrB2 HfB2 WC  1.  introduction  Carbon base materials such as graphite and carbon-carbon composites have several applications such as in aerospace, furnace elements, turbine blades and etc. These materials have suitable behavior at high temperatures (2500 °C) without need for cooling fluids such as high elastic modulus, high wear and fatigue resistance, high thermal shock resistance, low density and higher strength above 1200 °C [1-4]. The major drawback of these materials is their severe oxidation above 400 °C in the oxidized atmosphere [5]. In order to overcome this problem as well as protection against ablation, the ceramic coatings were used by various methods such as pack cementation, plasma spray, sol gel, electrolysis, chemical vapor deposition and etc. In general, the produced coatings by the above methods had a porous surface with small cracks as well as a limited thickness [6-14]. Ultra high temperature ceramics (UHTC) based on carbides and borides have a melting temperature above 3000 °C which are the best candidates for ablation-resistant coating on carbon components in aerospace applications [15,16]. The spark plasma sintering (SPS) is an effective method to produce materials that can be difficult to be sintered by conventional methods. The SPS can also be used for bonding of [17-19]. High heating rate and short dissimilar ceramics and metals process time are the main advantage of this method which leads to the formation of dense and fine microstructure [20]. ZrB2 based composites with SiC whiskers and various morphologies of carbon such and nano diamond [21], pulverized carbon fibers [22] and carbon nanoparticles [23] were sintered by SPS. In order to improve the ablation resistance of carbon-carbon composite, the ZrB2-SiC coating was used by powder pack cementation method with the thickness of 100-150 μm and dense microstructure without any macro-cracks or porosity. A number of micro-cracks were created during cooling to room temperature due to thermal expansion coefficients mismatch of substrate and coating [24]. In another research Oxidation resistance of carbon-carbon composite was improved with SiC (bond coat)/ZrB2-SiC-Si composite coating by the powder pack cementation. Their results showed that the coated sample had better oxidation resistance [25]. ZrB2-SiC double thick layer coating with a dense microstructure was produced by pressure diffusion. Results showed that this coating provides desirable oxidation protection for graphite and carbon-carbon composite under high heat flux and high temperature [26]. In another study, ZrB2-SiC-WC composite coating was produced by SPS process which had dense microstructure, no crack, and adequate adhesion to the diffusive SiC coating (sub-layer) [27]. On the knowledge of the authors, there is no report on the spark plasma sintering of ZrB2-SiC-HfB2-WC composite coating and its ablation properties in the literature. The aim of this study is to improve the ablation resistance of graphite substrate by UHTC coating. For the first time, the ZrB2-SiC composite coating with HfB2-WC additives was manufactured by SPS method. The influence of these additives was  ∗ Corresponding author. ∗∗ Corresponding author. E-mail addresses: m.sobhani@semnan.ac.ir (M. Sobhani), m_zakeri@merc.ac.ir (M. Zakeri).  https://doi.org/10.1016/j.ceramint.2020.06.297 Received 23 March 2020; Received in revised form 25 June 2020; Accepted 27 June 2020  Available online 01 July 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', \"S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  Table 1  Chemical composition of the manufactured coatings by SPS.  Sample  ZrB2 (V%)  SiC (V%)  Si (V%)  HfB2 (V%)  WC (V%)  A B C D E  70 15 15 66 14.25 14.25 64.75 13.875 13.875 63 13.5 13.5 Graphite substrate without coating  0 2.5 3.75 5  0 2.5 3.75 5  investigated on the ablation properties of the coating.  2. Experimental  Graphite discs (Ø30 × 5 mm) with a density of 1.73 g cm−3 were used as substrate. All disks were polished and washed in ethanol with ultrasonic in order to clean their surface and then dried at 120 °C for 2 h ZrB2 (D50 of 3 μm, 99.9%), SiC (D50 of 10 μm, 99%), Si (D50 of 15 μm, 99%), HfB2 (D50 of 5 μm, 99.9%) and WC (D50 of 10 μm, 99%) powders were used as starting materials to fabricate the composite coatings. The powders of each coating were weighed according to Table 1. These powders were mixed in an ultrasonic ethanol bath for 45 min. The obtained suspension was stirred and dried for 2 h with a magnet device and finally were completely dried at 150 °C for 2 h. The all inner surfaces of the mold were covered with graphite foil (Thickness, 1 mm) to prevent the powder contact to the graphite mold. The Coatings were applied on the graphite substrate in a graphite mold with a diameter of 30 mm at 1950 °C under uniaxial pressure 25 MPa for 30 min in a vacuum of 15 Pa by the SPS instrument of Easy Fashion Co. (SPS-20T-10, China) which was schematically shown in Fig. 1. Microstructures of the coated specimens were examined by field emission scanning electron microscopy (FESEM, MIRA3, TESCAN Co. Czech Republic) equipped with energy dispersive spectroscopy (EDS). The crystalline phases were identified by X-ray diffraction (XRD, PW3710, Philips Co. Netherlands). The ablations of the coatings were evaluated by oxyacetylene flame test with acetylene (98%) and oxygen (99.2%) gases. The heat flux, sample distance to burner, and test time loss (ΔW) and were 8500 w.m2, 15 mm and 60S, respectively. Weight ablation rate (R) were calculated on the basis of following formulas:  % ΔW  =  m  0  −  m  0  m  0  R  =  m  1  −  Δt  m  1  100  ∗  (1)  (2)  m0 and m1 are the sample weights before and after spectively. Δt is the ablation test time.  the ablation,  re Fig. 1. Schematic of coating preparation using spark plasma sintering.  Fig. 2. XRD patterns from the surface of the coated specimens.  3. Results and discussion  3.1. Physical properties  XRD patterns of the coated specimens with SPS were shown in Fig. 2. All of the peaks belong to the starting materials and no change was occurred in the primary phases except the Si. As can be seen, the Si phase was not identified in all samples due to its consumption during the SPS process. The starting Si penetrates to the porosities of the graphite substrate and forms secondary SiC with the carbon of substrate which cannot be distinguished from the primary SiC in the starting materials. These observations indicate the formation of a multi-phase composite coating on the graphite substrate that is expected to have better ablation resistance than the uncoated sample. The carbon peaks which were observed in some samples related to the graphite foil attached to the sample surface from SPS mold. WC reflections can clearly be seen in all samples except sample A (without additive). It means that no reaction was occurred between WC and other constituents. Zirconium diboride and hafnium diboride reflections were completely overlapped due to their lattice similarity and cannot be separated in these patterns. The penetration of molten Si and formation of SiC fills the coating-substrate interface as well as the spaces between other constituents and the porosities of the substrate. In the other words, this secondary SiC acts as a bonding agent to bind the composite's constituents together as well as to the substrate which leads to formation of a dense composite coating and prevents the penetration of oxygen into the substrate. The cross-section SEM images of sample C as well as their EDS were shown in Fig. 3. As seen, various phases can be distinguished on the basis of color contrast because of their different atomic mass and density. EDS analysis and elemental composition show that the points A, B, C, D, E and F are ZrB2, SiC, ZrB2, WC, HfB2 and ZrO2, respectively. Points A and C are ZrB2 phase with gray color. As discussed above, ZrB2 and HfB2 have very similar lattice which can easily form solid solution. These Points may be differ in their solved HfB2 amount in the ZrB2 lattice. Point E is HfB2 with white color due to its high atomic mass (200.11 g mol−1) and density (10.5 g cm−3) in compare with other constituents except WC with the atomic mass of 195.85 g mol−1 and density of 15.63 g cm−3. Points D and E have approximately same colors due to their similar atomic mass and density, but their elemental analysis confirm the composition of point D as WC. Point B is SiC with the dark color because of its low atomic mass (40.1 g mol−1) and density (3.21 g cm−3). Point F has similar composition to points A and C with some oxide layer on its surface which may be introduced from the starting materials or may be oxidized during preparation or sintering process.  25107  \\x0c\", 'S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  Fig. 3. SEM images of the cross-section microstructure of sample C with EDS analysis.  25108  \\x0c', \"S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  Fig. 4. Cross-section SEM image of sample C; a) coating thickness, b) Si penetration depth, c) coating and substrate and D) Si map analysis of C.  Table 2  Results of ablation including weight loss (ΔW) and rate of ablation (R).  Sample  A B C D E  ΔW (%)  4.8 4.4 2.0 2.7 7.4  R (mg.s-1)  7.3 6.5 3.0 4.2 7.8  Fig. 5. XRD pattern from the surface of the coating after ablation.  The coating thickness and it's the penetration depth in substrate were shown in Fig. 4. All samples had very similar microstructure and thickness. Therefore the SEM image of sample C was only shown and discussed here. All samples had similar thickness about 500 μm (Fig. 4A) and the penetration depth about 250 μm (Fig. 4-B). Graphite substrate and coating was strongly interconnected and no separation was observed in their interface. In the samples with additives, denser coating with lower porosity and deeper penetration was obtained. Penetration of the molten Si in the graphite porosities and their chemical reaction leads to the formation of SiC during SPS. These phenomenons were significantly accelerated by the applied pressure of SPS which leads to the improvement of mechanical and ablation properties of the coatings [27]. Fig. 4-D shows the Si map analysis of coating and substrate from Fig. 4-C. During the SPS process, as the temperature rises above the melting point of Si, the melted Si penetrates into the porosities of graphite substrate due to the capillary force and SPS pressure which can clearly be seen in Fig. 4-D. SiC grains were dominantly located in the inner part of the coating and substrate. The presence of SiC phase in the porosities of the graphite substrate increases its ablation resistance and adhesive strength in compare with pure graphite [28]. Homogeneity and penetration depth of Si were dependent on the several factors such as; graphite porosity and its size and distribution, SPS temperature, pressure and time which can influence the final properties of coating. In general it can be said that higher adhesive strength and ablation resistance will be obtained with more  25109  \\x0c\", 'S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  Fig. 6. Cross-section SEM image of the samples: A) Sample A without additive, B) sample B with 2.5%, C) sample C with 3.75% and D) sample D with 5% of each additives and also the EDS analyses of squires 1 and 2 on Fig. 6-C after ablation.  and uniform penetration at higher SPS temperature and pressure.  3.2. Ablation properties  The ablation resistances of the prepared coatings were investigated by high temperature oxyacetylene flam test. The results of this test including weight loss (%ΔW) and rate of ablation (R) were presented in Table 2. As can be seen, the ablation resistances of the coated samples  were higher than the uncoated sample. In fact, the uncoated sample was severely ablated. Sample C had the maximum ablation resistance in compare with the other coated samples, which may be due to the good match of the thermal expansion coefficient of the coating and its substrate. The XRD pattern of the coated specimen after oxyacetylene test at 2050 °C was shown in Fig. 5. All reflections in this pattern are related to ZrO2 and HfO2 due to the oxidation of ZrB2 and HfB2. Other oxidized constituents may be evaporated during ablation test. Ablation of the  25110  \\x0c', 'S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  coated samples may be performed on the basis of  following reactions:  ZrB2(s)+5/2O2(g) = ZrO2(s)+B2O3(g)  HfB2(s)+5/2O2(g) = HfO2(s)+B2O3(g)  SiC(s)+2O2(g) = SiO2(g)+CO2(g)  SiO2(s) = SiO(g)+1/2O2(g)  WC(s)+5/2O2(g) = WO3(g)+CO2(g)  (3)  (4)  (5)  (6)  (7)  In the first stage of flame test, oxidizing of the composite constituents leads to the formation of some oxides on the basis of above reactions. These oxides form a viscose silica base glass which adheres to the surface of the coating and protects the sub layers and graphite substrate from more oxidation. In the next stage, evaporation of SiO and WO3 leads to the formation of porous microstructure in the coating. It means that these phenomena are the main mechanisms of oxidation/ ablation resistance of coating. SEM images of the samples after the flame test were shown in Fig. 6. As can be seen, sample A (Fig. 6-A) has a dense surface without any porosity and only severe degradation (ablation) was occurred on its surface. In sample B (Fig. 6-B) with 2.5% of each additive, its ablation was occurred from the surface as well as in the coating (oxidized thickness approx. of 150 μm). the inner parts of On the basis of above reactions, the oxidation gaseous products led to the formation of porosity in the coating. Some surface cracking can be seen in the microstructure of ablated coating due to the thermal expansion coefficient mismatch of the composite constituents, ablation products and the substrate. No surface degradation (ablation) was occurred on the surface of the sample C with 3.75% of each additive (Fig. 6-C). On the other hand, bulk ablation was occurred by the oxidation of additives (WC, SiC and HfB2) which led to the formation of porous microstructure (oxidized thickness approx. of 250 μm). As seen, the main structure (ZrB2) of the coating was not changed in this sample. This sample had the minimum ablation rate (Table 2) which may be due to the higher oxidation rate of WC or higher thermal conductivity of HfB2 in compare with other constituents. The energy of flame is consumed during oxidation of WC and/or dissipated in the body of coating and substrate which leads to the decreasing of overall temperature and ablation rate. Sample D (Fig. 6-D) was similar to the sample C as discussed before (oxidized thickness approx. of 200 μm). It is clear from Fig. 6 that ablation is a surface parameter. It means that ablation is initiated from the surface of the sample and gradually propagated to reach the graphite sub layer. Therefore the ablation of the coating increases the working time of the graphite substrate in its application. But there is difference between samples C with others. In this sample, coating saves its main structure and the formed porous section in the microstructure act as thermal barrier and insulation which decreases the sub layers and substrate temperature and increases the ablation resistance and working time of the specimen. The EDS analyses of sample C in the ablated and not ablated sections of the coating were presented in Fig. 6. As seen in the ablated section (Squire 1Fig. 6-F), oxygen content is in its maximum state and Si and W contents are in their minimum state. On the other hand, in the non-ablated section (Squire 2Fig. 6-E), oxygen could not diffuse to this region and Si and W contents are higher than squire 1 which confirms the above discussed ablation mechanisms.  4. Conclusion  The ZrB2-SiC composite coatings with HfB2 and WC additives were successfully manufactured on the graphite substrate using SPS. These coatings significantly improved the ablation resistance of graphite. Sample C with the composition of 3.75% of each additive exhibited the maximum ablation resistance compared to other samples. In fact, evaporation of gaseous products during the flame test, lead to the consumption of flame energy and remaining of composite main structure  (ZrB2) of coating for a longer periods of time. Higher thermal conductivity of the composite because of its higher amount of HfB2 leads to the dissipating of flame energy, decreasing of surface overall temperature and increasing of ablation resistance.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  References  [1]  [8]  [9]  [5]  [6]  [2]  [3]  [14]  J.D. Webster, M.E. Westwood, F.H. Hayes, R.J. Day, R. Taylor, A. Duran, et al., Oxidation protection coatings for C/SiC based on yttrium silicate, J. Eur. Ceram. Soc. 18 (16) (1998) 2345-2350. E. Fitzer, L.M. Manocha, Carbon Reinforcements and Carbon/carbon Composites, Springer, Berlin, 1998. J.R. Strife, J.E. Sheehan, Ceramic coatings for carbon-carbon composites, Am. Ceram. Soc. Bull. 67 (2) (1988) 369-374. [4] D.E. Wittmer, M.Z. Temuri, Thermochemical studies in selected metal-carbonoxygen systems, J. Am. Ceram. Soc. 74 (5) (1991) 973-982. J. Chown, R.F. Deacon, N. Singer, A.E.S. White, Refractory cpatings on graphite, Special Ceram. 81 (1963). J.-H. Kang, H.-M. Cha, J.-H. Lee, Y.-G. Jung, H.-S. Lee, V. Srivastava, Synthesis and characterization of non-oxide ceramic coatings on graphite substrates by solid-vapor reaction process, Prog. Org. Coating 61 (2008) 291-299. [7] H. Jian-Feng, Z. Xie-Rong, L. He-Jun, X. Xin-Bo, F. Ye-wei, Influence of the preparation temperature on the phase, microstructure and anti-oxidation property of a SiC coating for C/C composites, Carbon 42 (2004) 1517-1521. F. Qian-Gang, L. He-Jun, W. Yong-Jie, L. Ke-Zhi, W. Heng, A Si-SiC oxidation protective coating for carbon/carbon composites prepared by a two-step pack ceInt. 35 (2009) 2525-2529. mentation, Ceram. C. Hu, Y. Niu, H. Li, M. Ren, X. Zheng, J. Sun, SiC coatings for carbon/carbon composites fabricated by vacuum plasma spraying technology, J. Therm. Spray Technol. 21 (2012) 16-22. [10] Q.G. Fu, X. Hui, H.-J. Li, K.-Z. Li, X.H. Shi, Z. Hua, Anti-oxidation property of a multi-layer coating for carbon/carbon composites in a wind tunnel at 1500 °C, N. Carbon Mater. 25 (2010) 279-284. [11] K.-T. Wang, L.-Y. Cao, J.-F. Huang, J. Fei, A mullite/SiC oxidation protective coating for carbon-carbon composites, J. Eur. Ceram. Soc. 33 (2013) 191-198. [12] X. Ren, H. Li, Q. Fu, K. Li, TaxHf1− xB2-SiC multiphase oxidation protective coating for SiC-coated carbon-carbon composites, Corrosion Sci. 87 (2014) 479-488. [13] X.Y. Yao, H.J. Li, Y.L. Zhang, J.J. Ren, D.J. Yao, J. Tao, A SiC/ZrB2-SiC/SiC oxidation resistance multilayer coating for carbon/carbon composites, Corrosion Sci. 57 (2012) 148-153. Y. Zhang, H. Hu, J. Ren, T. Li, T. Fei, C. Wang, Effect of the surface microstructure of SiC inner coating on the bonding strength and ablation resistance of ZrB2-SiC Int. 42 (2016) 18657-18665. coating for C/C composites, Ceram. L. Silvestroni, H.J. Kleebe, W.G. Fahrenholtz, J. Watts, Super-strong materials for temperatures exceeding 2000 °C, Sci. Rep. 7 (2017) 1-8. F. Monteverde, L. Scatteia, Resistance to thermal shock and to oxidation of metal Diborides-SiC ceramics for aerospace application, J. Am. Ceram. Soc. 90 (4) (2007) 1130-1138. T. Okuni, Y. Miyamoto, H. Abe, M. Naito, Joining of silicon carbide and graphite by spark plasma sintering, Ceram. Int. 40 (2014) 1359-1363. [18] A. Singh, S.R. Bakshi, D.A. Virzi, A.K. Keshri, A. Agarwal, S.P. Harimkar, In-situ synthesis of TiC/SiC/Ti 3 SiC 2 composite coatings by spark plasma sintering, Surf. Coating. Technol. 205 (2011) 3840-3846. Y. Jiang, J. Yang, Z. Xie, R. Gao, Q. Fang, Corrosion resistance of W-Cr-C coatings fabricated by spark plasma sintering method, Surf. Coating. Technol. 254 (2014) 202-206. [20] D. Ghahremani, T. Ebadzadeh, A. Maghsodipour, Spark plasma sintering of mullite: relation between microstructure, properties and spark plasma sintering (SPS) Int. 41 (2015) 6409-6416. parameters, Ceram. [21] Mehdi Fattahi, Yashar Azizian-Kalandaragh, Seyed Ali Delbari, Abbas Sabahi Naminie, Zohre Ahmadi, Mehdi Shahedi Asl, Nano-diamond reinforced ZrB2-SiC composites, Ceram. Int. 46 (8) (2020) 10172-10179 Part A. [22] Mehdi Shahedi Asl, Yashar Azizian-Kalandaragh, Zohre Ahmadi, Abbas Sabahi Naminie, Amir Motallebzadeh, Spark plasma sintering of ZrB2-based composites co-reinforced with SiC whiskers and pulverized carbon fibers, Int. J. Refract. Metals Hard Mater. 83 (September) (2019) 104989. [23] Azizian-Kalandaragh Yashar, Abbas Sabahi Namini, Zohre Ahmadi, Mehdi Shahedi Asl, Reinforcing effects of SiC whiskers and carbon nanoparticles in spark plasma sintered ZrB2 matrix composites, Ceram. Int. 44 (16) (November 2018) 19932-19938. [24] X. Zou, Q. Fu, L. Liu, H. Li, Y. Wang, X. Yao, Z. He, ZrB2-SiC coating to protect carbon/carbon composites against ablation, Surf. Coating. Technol. 226 (2013) 17-21. [25] X. Yao, H. Li, Y. i Zhang, H. Wu, X. Qiang, A SiC-Si-ZrB2 multiphase oxidation protective ceramic coating for SiC-coated carbon/carbon composites, Ceram. Int. 38  [15]  [16]  [17]  [19]  25111  \\x0c', 'S.A. Akbarpour Shalmani, et al.  Ceramics International 46 (2020) 25106-25112  [26]  (2012) 2095-2100. E.L. Corral, R.E. Loehman, Ultra-high-temperature ceramic coatings for oxidation protection of carbon-carbon composites, J. Am. Ceram. Soc. 91 (5) (2008) 1495-1502. [27] M. Shirani, M. Rahimipour, M. Zakeri, S. Safi, T. Ebadzadeh, ZrB2-SiC-WC coating  with SiC diffusion bond coat on graphite by spark plasma sintering process, Ceram.  Int. 43 (2017) 14517-14520. S. Safi, A. Kazemzadeh, MCMB-SiC composites; new class high-temperature structural materials for aerospace applications, Ceram. Int. 39 (2013) 81-86.  [28]  25112  \\x0c']"
},{
  "_id": 43,
  "PDF": "Effect of HfB2and WC additives on the ablation resistance of ZrB2–SiC composite coating manufactured by SPS.pdf",
  "Text": "['Ceramics International xxx (xxxx) xxx-xxx  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Effect of HfB2 and WC additives on the ablation resistance of ZrB2-SiC composite coating manufactured by SPS  S. Ali Akbarpour Shalmania, M. Sobhania,∗, O. Mirzaeea, M. Zakerib,∗∗  a Faculty of Materials and Metallurgical Engineering, Semnan University, Semnan, b Ceramic Dept, Materials and Energy Research Center, Karaj, Iran  Iran  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ablation Composite SPS ZrB2 HfB2 WC  1.  introduction  ZrB2-SiC composite coating with HfB2-WC as additives was successfully manufactured on graphite substrate by spark plasma sintering method. The effect of HfB2 and WC on coating properties was investigated. The microstructure, phase changes and ablation resistance of the coatings were studied. Results showed that composite coatings with the thickness of 500 μm and penetration depth of 250 μm were obtained for all samples. Results of the oxyacetylene flame test showed that the ablation resistance of the coated graphite was significantly increased. The minimum ablation weight percentage and rate of 2% and 3 mg s−1 were obtained for the composite coating with 3.75% of each additive.  Carbon base materials such as graphite and carbon-carbon composites have several applications such as in aerospace, furnace elements, turbine blades and etc. These materials have suitable behavior at high temperatures (2500 °C) without need for cooling fluids such as high elastic modulus, high wear and fatigue resistance, high thermal shock resistance, low density and higher strength above 1200 °C [1-4]. The major drawback of these materials is their severe oxidation above 400 °C in the oxidized atmosphere [5]. In order to overcome this problem as well as protection against ablation, the ceramic coatings were used by various methods such as pack cementation, plasma spray, sol gel, electrolysis, chemical vapor deposition and etc. In general, the produced coatings by the above methods had a porous surface with small cracks as well as a limited thickness [6-14]. Ultra high temperature ceramics (UHTC) based on carbides and borides have a melting temperature above 3000 °C which are the best candidates for ablation-resistant coating on carbon components in aerospace applications [15,16]. The spark plasma sintering (SPS) is an effective method to produce materials that can be difficult to be sintered by conventional methods. The SPS can also be used for bonding of [17-19]. High heating rate and short dissimilar ceramics and metals process time are the main advantage of this method which leads to the formation of dense and fine microstructure [20]. ZrB2 based composites with SiC whiskers and various morphologies of carbon such and nano diamond [21], pulverized carbon fibers [22] and carbon nanoparticles [23] were sintered by SPS. In order to improve the ablation resistance of carbon-carbon composite, the ZrB2-SiC coating was used by powder pack cementation method with the thickness of 100-150 μm and dense microstructure without any macro-cracks or porosity. A number of micro-cracks were created during cooling to room temperature due to thermal expansion coefficients mismatch of substrate and coating [24]. In another research Oxidation resistance of carbon-carbon composite was improved with SiC (bond coat)/ZrB2-SiC-Si composite coating by the powder pack cementation. Their results showed that the coated sample had better oxidation resistance [25]. ZrB2-SiC double thick layer coating with a dense microstructure was produced by pressure diffusion. Results showed that this coating provides desirable oxidation protection for graphite and carbon-carbon composite under high heat flux and high temperature [26]. In another study, ZrB2-SiC-WC composite coating was produced by SPS process which had dense microstructure, no crack, and adequate adhesion to the diffusive SiC coating (sub-layer) [27]. On the knowledge of the authors, there is no report on the spark plasma sintering of ZrB2-SiC-HfB2-WC composite coating and its ablation properties in the literature. The aim of this study is to improve the ablation resistance of graphite substrate by UHTC coating. For the first time, the ZrB2-SiC composite coating with HfB2-WC additives was manufactured by SPS method. The influence of these additives was  ∗ Corresponding author. ∗∗ Corresponding author. E-mail addresses: m.sobhani@semnan.ac.ir (M. Sobhani), m_zakeri@merc.ac.ir (M. Zakeri).  https://doi.org/10.1016/j.ceramint.2020.06.297 Received 23 March 2020; Received in revised form 25 June 2020; Accepted 27 June 2020  0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  Please cite this article as: S. Ali Akbarpour Shalmani, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2020.06.297  \\x0c', \"S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  Table 1  Chemical composition of the manufactured coatings by SPS.  Sample  ZrB2 (V%)  SiC (V%)  Si (V%)  HfB2 (V%)  WC (V%)  A B C D E  70 15 15 66 14.25 14.25 64.75 13.875 13.875 63 13.5 13.5 Graphite substrate without coating  0 2.5 3.75 5  0 2.5 3.75 5  investigated on the ablation properties of the coating.  2. Experimental  Graphite discs (Ø30 × 5 mm) with a density of 1.73 g cm−3 were used as substrate. All disks were polished and washed in ethanol with ultrasonic in order to clean their surface and then dried at 120 °C for 2 h ZrB2 (D50 of 3 μm, 99.9%), SiC (D50 of 10 μm, 99%), Si (D50 of 15 μm, 99%), HfB2 (D50 of 5 μm, 99.9%) and WC (D50 of 10 μm, 99%) powders were used as starting materials to fabricate the composite coatings. The powders of each coating were weighed according to Table 1. These powders were mixed in an ultrasonic ethanol bath for 45 min. The obtained suspension was stirred and dried for 2 h with a magnet device and finally were completely dried at 150 °C for 2 h. The all inner surfaces of the mold were covered with graphite foil (Thickness, 1 mm) to prevent the powder contact to the graphite mold. The Coatings were applied on the graphite substrate in a graphite mold with a diameter of 30 mm at 1950 °C under uniaxial pressure 25 MPa for 30 min in a vacuum of 15 Pa by the SPS instrument of Easy Fashion Co. (SPS-20T-10, China) which was schematically shown in Fig. 1. Microstructures of the coated specimens were examined by field emission scanning electron microscopy (FESEM, MIRA3, TESCAN Co. Czech Republic) equipped with energy dispersive spectroscopy (EDS). The crystalline phases were identified by X-ray diffraction (XRD, PW3710, Philips Co. Netherlands). The ablations of the coatings were evaluated by oxyacetylene flame test with acetylene (98%) and oxygen (99.2%) gases. The heat flux, sample distance to burner, and test time loss (ΔW) and were 8500 w.m2, 15 mm and 60S, respectively. Weight ablation rate (R) were calculated on the basis of following formulas:  % ΔW  =  m  0  −  m  0  m  0  R  =  m  1  −  Δt  m  1  100  ∗  (1)  (2)  m0 and m1 are the sample weights before and after spectively. Δt is the ablation test time.  the ablation,  re Fig. 1. Schematic of coating preparation using spark plasma sintering.  2  Fig. 2. XRD patterns from the surface of the coated specimens.  3. Results and discussion  3.1. Physical properties  XRD patterns of the coated specimens with SPS were shown in Fig. 2. All of the peaks belong to the starting materials and no change was occurred in the primary phases except the Si. As can be seen, the Si phase was not identified in all samples due to its consumption during the SPS process. The starting Si penetrates to the porosities of the graphite substrate and forms secondary SiC with the carbon of substrate which cannot be distinguished from the primary SiC in the starting materials. These observations indicate the formation of a multi-phase composite coating on the graphite substrate that is expected to have better ablation resistance than the uncoated sample. The carbon peaks which were observed in some samples related to the graphite foil attached to the sample surface from SPS mold. WC reflections can clearly be seen in all samples except sample A (without additive). It means that no reaction was occurred between WC and other constituents. Zirconium diboride and hafnium diboride reflections were completely overlapped due to their lattice similarity and cannot be separated in these patterns. The penetration of molten Si and formation of SiC fills the coating-substrate interface as well as the spaces between other constituents and the porosities of the substrate. In the other words, this secondary SiC acts as a bonding agent to bind the composite's constituents together as well as to the substrate which leads to formation of a dense composite coating and prevents the penetration of oxygen into the substrate. The cross-section SEM images of sample C as well as their EDS were shown in Fig. 3. As seen, various phases can be distinguished on the basis of color contrast because of their different atomic mass and density. EDS analysis and elemental composition show that the points A, B, C, D, E and F are ZrB2, SiC, ZrB2, WC, HfB2 and ZrO2, respectively. Points A and C are ZrB2 phase with gray color. As discussed above, ZrB2 and HfB2 have very similar lattice which can easily form solid solution. These Points may be differ in their solved HfB2 amount in the ZrB2 lattice. Point E is HfB2 with white color due to its high atomic mass (200.11 g mol−1) and density (10.5 g cm−3) in compare with other constituents except WC with the atomic mass of 195.85 g mol−1 and density of 15.63 g cm−3. Points D and E have approximately same colors due to their similar atomic mass and density, but their elemental analysis confirm the composition of point D as WC. Point B is SiC with the dark color because of its low atomic mass (40.1 g mol−1) and density (3.21 g cm−3). Point F has similar composition to points A and C with some oxide layer on its surface which may be introduced from the starting materials or may be oxidized during preparation or sintering process.  \\x0c\", 'S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 3. SEM images of the cross-section microstructure of sample C with EDS analysis.  3  \\x0c', \"S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 4. Cross-section SEM image of sample C; a) coating thickness, b) Si penetration depth, c) coating and substrate and D) Si map analysis of C.  Table 2  Results of ablation including weight loss (ΔW) and rate of ablation (R).  Sample  A B C D E  ΔW (%)  4.8 4.4 2.0 2.7 7.4  R (mg.s-1)  7.3 6.5 3.0 4.2 7.8  Fig. 5. XRD pattern from the surface of the coating after ablation.  4  The coating thickness and it's the penetration depth in substrate were shown in Fig. 4. All samples had very similar microstructure and thickness. Therefore the SEM image of sample C was only shown and discussed here. All samples had similar thickness about 500 μm (Fig. 4A) and the penetration depth about 250 μm (Fig. 4-B). Graphite substrate and coating was strongly interconnected and no separation was observed in their interface. In the samples with additives, denser coating with lower porosity and deeper penetration was obtained. Penetration of the molten Si in the graphite porosities and their chemical reaction leads to the formation of SiC during SPS. These phenomenons were significantly accelerated by the applied pressure of SPS which leads to the improvement of mechanical and ablation properties of the coatings [27]. Fig. 4-D shows the Si map analysis of coating and substrate from Fig. 4-C. During the SPS process, as the temperature rises above the melting point of Si, the melted Si penetrates into the porosities of graphite substrate due to the capillary force and SPS pressure which can clearly be seen in Fig. 4-D. SiC grains were dominantly located in the inner part of the coating and substrate. The presence of SiC phase in the porosities of the graphite substrate increases its ablation resistance and adhesive strength in compare with pure graphite [28]. Homogeneity and penetration depth of Si were dependent on the several factors such as; graphite porosity and its size and distribution, SPS temperature, pressure and time which can influence the final properties of coating. In general it can be said that higher adhesive strength and ablation resistance will be obtained with more  \\x0c\", 'S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 6. Cross-section SEM image of the samples: A) Sample A without additive, B) sample B with 2.5%, C) sample C with 3.75% and D) sample D with 5% of each additives and also the EDS analyses of squires 1 and 2 on Fig. 6-C after ablation.  and uniform penetration at higher SPS temperature and pressure.  3.2. Ablation properties  The ablation resistances of the prepared coatings were investigated by high temperature oxyacetylene flam test. The results of this test including weight loss (%ΔW) and rate of ablation (R) were presented in Table 2. As can be seen, the ablation resistances of the coated samples  were higher than the uncoated sample. In fact, the uncoated sample was severely ablated. Sample C had the maximum ablation resistance in compare with the other coated samples, which may be due to the good match of the thermal expansion coefficient of the coating and its substrate. The XRD pattern of the coated specimen after oxyacetylene test at 2050 °C was shown in Fig. 5. All reflections in this pattern are related to ZrO2 and HfO2 due to the oxidation of ZrB2 and HfB2. Other oxidized constituents may be evaporated during ablation test. Ablation of the  5  \\x0c', 'S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  coated samples may be performed on the basis of  following reactions:  ZrB2(s)+5/2O2(g) = ZrO2(s)+B2O3(g)  HfB2(s)+5/2O2(g) = HfO2(s)+B2O3(g)  SiC(s)+2O2(g) = SiO2(g)+CO2(g)  SiO2(s) = SiO(g)+1/2O2(g)  WC(s)+5/2O2(g) = WO3(g)+CO2(g)  (3)  (4)  (5)  (6)  (7)  In the first stage of flame test, oxidizing of the composite constituents leads to the formation of some oxides on the basis of above reactions. These oxides form a viscose silica base glass which adheres to the surface of the coating and protects the sub layers and graphite substrate from more oxidation. In the next stage, evaporation of SiO and WO3 leads to the formation of porous microstructure in the coating. It means that these phenomena are the main mechanisms of oxidation/ ablation resistance of coating. SEM images of the samples after the flame test were shown in Fig. 6. As can be seen, sample A (Fig. 6-A) has a dense surface without any porosity and only severe degradation (ablation) was occurred on its surface. In sample B (Fig. 6-B) with 2.5% of each additive, its ablation was occurred from the surface as well as in the coating (oxidized thickness approx. of 150 μm). the inner parts of On the basis of above reactions, the oxidation gaseous products led to the formation of porosity in the coating. Some surface cracking can be seen in the microstructure of ablated coating due to the thermal expansion coefficient mismatch of the composite constituents, ablation products and the substrate. No surface degradation (ablation) was occurred on the surface of the sample C with 3.75% of each additive (Fig. 6-C). On the other hand, bulk ablation was occurred by the oxidation of additives (WC, SiC and HfB2) which led to the formation of porous microstructure (oxidized thickness approx. of 250 μm). As seen, the main structure (ZrB2) of the coating was not changed in this sample. This sample had the minimum ablation rate (Table 2) which may be due to the higher oxidation rate of WC or higher thermal conductivity of HfB2 in compare with other constituents. The energy of flame is consumed during oxidation of WC and/or dissipated in the body of coating and substrate which leads to the decreasing of overall temperature and ablation rate. Sample D (Fig. 6-D) was similar to the sample C as discussed before (oxidized thickness approx. of 200 μm). It is clear from Fig. 6 that ablation is a surface parameter. It means that ablation is initiated from the surface of the sample and gradually propagated to reach the graphite sub layer. Therefore the ablation of the coating increases the working time of the graphite substrate in its application. But there is difference between samples C with others. In this sample, coating saves its main structure and the formed porous section in the microstructure act as thermal barrier and insulation which decreases the sub layers and substrate temperature and increases the ablation resistance and working time of the specimen. The EDS analyses of sample C in the ablated and not ablated sections of the coating were presented in Fig. 6. As seen in the ablated section (Squire 1Fig. 6-F), oxygen content is in its maximum state and Si and W contents are in their minimum state. On the other hand, in the non-ablated section (Squire 2Fig. 6-E), oxygen could not diffuse to this region and Si and W contents are higher than squire 1 which confirms the above discussed ablation mechanisms.  4. Conclusion  The ZrB2-SiC composite coatings with HfB2 and WC additives were successfully manufactured on the graphite substrate using SPS. These coatings significantly improved the ablation resistance of graphite. Sample C with the composition of 3.75% of each additive exhibited the maximum ablation resistance compared to other samples. In fact, evaporation of gaseous products during the flame test, lead to the consumption of flame energy and remaining of composite main structure  6  (ZrB2) of coating for a longer periods of time. Higher thermal conductivity of the composite because of its higher amount of HfB2 leads to the dissipating of flame energy, decreasing of surface overall temperature and increasing of ablation resistance.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  References  [1]  [8]  [9]  [5]  [6]  [2]  [3]  [14]  J.D. Webster, M.E. Westwood, F.H. Hayes, R.J. Day, R. Taylor, A. Duran, et al., Oxidation protection coatings for C/SiC based on yttrium silicate, J. Eur. Ceram. Soc. 18 (16) (1998) 2345-2350. E. Fitzer, L.M. Manocha, Carbon Reinforcements and Carbon/carbon Composites, Springer, Berlin, 1998. J.R. Strife, J.E. Sheehan, Ceramic coatings for carbon-carbon composites, Am. Ceram. Soc. Bull. 67 (2) (1988) 369-374. [4] D.E. Wittmer, M.Z. Temuri, Thermochemical studies in selected metal-carbonoxygen systems, J. Am. Ceram. Soc. 74 (5) (1991) 973-982. J. Chown, R.F. Deacon, N. Singer, A.E.S. White, Refractory cpatings on graphite, Special Ceram. 81 (1963). J.-H. Kang, H.-M. Cha, J.-H. Lee, Y.-G. Jung, H.-S. Lee, V. Srivastava, Synthesis and characterization of non-oxide ceramic coatings on graphite substrates by solid-vapor reaction process, Prog. Org. Coating 61 (2008) 291-299. [7] H. Jian-Feng, Z. Xie-Rong, L. He-Jun, X. Xin-Bo, F. Ye-wei, Influence of the preparation temperature on the phase, microstructure and anti-oxidation property of a SiC coating for C/C composites, Carbon 42 (2004) 1517-1521. F. Qian-Gang, L. He-Jun, W. Yong-Jie, L. Ke-Zhi, W. Heng, A Si-SiC oxidation protective coating for carbon/carbon composites prepared by a two-step pack ceInt. 35 (2009) 2525-2529. mentation, Ceram. C. Hu, Y. Niu, H. Li, M. Ren, X. Zheng, J. Sun, SiC coatings for carbon/carbon composites fabricated by vacuum plasma spraying technology, J. Therm. Spray Technol. 21 (2012) 16-22. [10] Q.G. Fu, X. Hui, H.-J. Li, K.-Z. Li, X.H. Shi, Z. Hua, Anti-oxidation property of a multi-layer coating for carbon/carbon composites in a wind tunnel at 1500 °C, N. Carbon Mater. 25 (2010) 279-284. [11] K.-T. Wang, L.-Y. Cao, J.-F. Huang, J. Fei, A mullite/SiC oxidation protective coating for carbon-carbon composites, J. Eur. Ceram. Soc. 33 (2013) 191-198. [12] X. Ren, H. Li, Q. Fu, K. Li, TaxHf1− xB2-SiC multiphase oxidation protective coating for SiC-coated carbon-carbon composites, Corrosion Sci. 87 (2014) 479-488. [13] X.Y. Yao, H.J. Li, Y.L. Zhang, J.J. Ren, D.J. Yao, J. Tao, A SiC/ZrB2-SiC/SiC oxidation resistance multilayer coating for carbon/carbon composites, Corrosion Sci. 57 (2012) 148-153. Y. Zhang, H. Hu, J. Ren, T. Li, T. Fei, C. Wang, Effect of the surface microstructure of SiC inner coating on the bonding strength and ablation resistance of ZrB2-SiC Int. 42 (2016) 18657-18665. coating for C/C composites, Ceram. L. Silvestroni, H.J. Kleebe, W.G. Fahrenholtz, J. Watts, Super-strong materials for temperatures exceeding 2000 °C, Sci. Rep. 7 (2017) 1-8. F. Monteverde, L. Scatteia, Resistance to thermal shock and to oxidation of metal Diborides-SiC ceramics for aerospace application, J. Am. Ceram. Soc. 90 (4) (2007) 1130-1138. T. Okuni, Y. Miyamoto, H. Abe, M. Naito, Joining of silicon carbide and graphite by spark plasma sintering, Ceram. Int. 40 (2014) 1359-1363. [18] A. Singh, S.R. Bakshi, D.A. Virzi, A.K. Keshri, A. Agarwal, S.P. Harimkar, In-situ synthesis of TiC/SiC/Ti 3 SiC 2 composite coatings by spark plasma sintering, Surf. Coating. Technol. 205 (2011) 3840-3846. Y. Jiang, J. Yang, Z. Xie, R. Gao, Q. Fang, Corrosion resistance of W-Cr-C coatings fabricated by spark plasma sintering method, Surf. Coating. Technol. 254 (2014) 202-206. [20] D. Ghahremani, T. Ebadzadeh, A. Maghsodipour, Spark plasma sintering of mullite: relation between microstructure, properties and spark plasma sintering (SPS) Int. 41 (2015) 6409-6416. parameters, Ceram. [21] Mehdi Fattahi, Yashar Azizian-Kalandaragh, Seyed Ali Delbari, Abbas Sabahi Naminie, Zohre Ahmadi, Mehdi Shahedi Asl, Nano-diamond reinforced ZrB2-SiC composites, Ceram. Int. 46 (8) (2020) 10172-10179 Part A. [22] Mehdi Shahedi Asl, Yashar Azizian-Kalandaragh, Zohre Ahmadi, Abbas Sabahi Naminie, Amir Motallebzadeh, Spark plasma sintering of ZrB2-based composites co-reinforced with SiC whiskers and pulverized carbon fibers, Int. J. Refract. Metals Hard Mater. 83 (September) (2019) 104989. [23] Azizian-Kalandaragh Yashar, Abbas Sabahi Namini, Zohre Ahmadi, Mehdi Shahedi Asl, Reinforcing effects of SiC whiskers and carbon nanoparticles in spark plasma sintered ZrB2 matrix composites, Ceram. Int. 44 (16) (November 2018) 19932-19938. [24] X. Zou, Q. Fu, L. Liu, H. Li, Y. Wang, X. Yao, Z. He, ZrB2-SiC coating to protect carbon/carbon composites against ablation, Surf. Coating. Technol. 226 (2013) 17-21. [25] X. Yao, H. Li, Y. i Zhang, H. Wu, X. Qiang, A SiC-Si-ZrB2 multiphase oxidation protective ceramic coating for SiC-coated carbon/carbon composites, Ceram. Int. 38  [15]  [16]  [17]  [19]  \\x0c', 'S.A. Akbarpour Shalmani, et al.  Ceramics International xxx (xxxx) xxx-xxx  [26]  (2012) 2095-2100. E.L. Corral, R.E. Loehman, Ultra-high-temperature ceramic coatings for oxidation protection of carbon-carbon composites, J. Am. Ceram. Soc. 91 (5) (2008) 1495-1502. [27] M. Shirani, M. Rahimipour, M. Zakeri, S. Safi, T. Ebadzadeh, ZrB2-SiC-WC coating  with SiC diffusion bond coat on graphite by spark plasma sintering process, Ceram.  Int. 43 (2017) 14517-14520. S. Safi, A. Kazemzadeh, MCMB-SiC composites; new class high-temperature structural materials for aerospace applications, Ceram. Int. 39 (2013) 81-86.  [28]  7  \\x0c']"
},{
  "_id": 44,
  "PDF": "Effect of La2O3 addition on long-term oxidation kinetics of ZrB2–SiC and HfB2–SiC ultra-high temperature ceramics.pdf",
  "Text": "['Available online at www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 34 (2014) 3535-3548  Effect of La2O3 addition on long-term oxidation kinetics of ZrB2-SiC and HfB2-SiC ultra-high temperature ceramics  E. Zapata-Solvas a,∗  , D.D. Jayaseelan a , P.M. Brown b , W.E. Lee a  a Centre for Advanced Structural Ceramics, Imperial College London, SW7 2AZ, UK b Dstl, Porton Down, Salisbury, Wiltshire SP4 0JQ, UK  Received 4 May 2014; received in revised form 28 May 2014; accepted 2 June 2014  Available online 1 July 2014  Abstract  Long-term oxidation kinetics of SiC-reinforced UHTCs and La2O3 -doped UHTCs over an intermediate temperature range (1400-1600 C) reveal partially protective behavior for the former characterized by an oxidation kinetic exponent 1 < n < 2. In addition, unstable oxidation behavior was observed in HfB2 -based UHTCs due to the presence of SiC agglomerates. On the other hand, La2O3 -doped UHTCs were found to be protective over the whole temperature range studied (n = 2), in particular at 1600 C, where oxidation kinetic exponents as high as 8 were observed as a consequence of formation of new oxidation protective particles, MeOxCy , where Me is Zr, Hf or Si. Adsorption of oxygen-containing species formed protective MeOxCy phases, which enhanced the thermal stability of the oxide scale as well as providing protection against oxidation for long exposure times at 1600 C. © 2014 Elsevier Ltd. All rights reserved.           Keywords: Oxidation kinetics; Oxidation resistance; Borosilicate coating; Protective coating; Coating instability  1.  Introduction     Ceramic materials with melting points in excess of 3000 C, so-called ultra-high temperature ceramics (UHTCs) such as zirconium and hafnium diborides and carbides,1 have potential for use in thermal protection systems (TPS) or refractory applications. In addition, zirconium and hafnium diborides have higher thermal conductivity (100 W/(m K) at room temC)2 perature and >50 W/(m K) at 1900 than their respective C).3 carbides (<40 W/(m K) from room temperature to 800 The combination of high temperature capability and high thermal conductivity is particularly beneficial and highly desirable for hypersonic applications, providing thermal transport by conduction and radiation during exposure to extremely        ∗  Corresponding author. Current address: Instituto de Ciencia de Materiales de  Sevilla (CSIC-Universidad de Sevilla), C/Américo Vespucio 49, 41092 Seville,  Spain. Tel.: +34 954 559978; fax: +34 954552870.  E-mail addresses: ezapata@us.es, ezapata@icmse.csic.es  (E. Zapata-Solvas).  http://dx.doi.org/10.1016/j.jeurceramsoc.2014.06.004  0955-2219/© 2014 Elsevier Ltd. All rights reserved.     high-temperature reactive environments.4 TPS for hypersonic applications include sharp nose cones (SNC) and sharp-leading edges (SLE) which improve maneuverability and maximum operating speed of hypersonic vehicles, opening up a new range of hypersonic and re-entry trajectories compared to aerospace vehicles equipped with blunt TPS such as the NASA space shuttle orbiter. However, SNC and SLE maximum surface temperatures increase slightly when the surface radius of curvature decreases under the same heat flux conditions5 (30 C difference at 1800 C from blunt, 5 mm, to SNC, 0.5 mm) and temperature reached during hypersonic re-entries range from 1500 C to 2000 C or even higher depending on re-entry conditions which controls the exposed heat flux.6,7 In addition, it is highly desirable for a hypersonic vehicle to be reusable and during the last 30 years the only reusable space vehicle has been the space shuttle which is equipped with blunt nose cone and leading edges. Therefore, long-term oxidation kinetic studies are required to understand how UHTCs behave for long exposure times. As mentioned above, ZrB2 and HfB2 potentially fulfill the stringent requirements for TPS in hypersonic applications. However, ZrB2           \\x0c', '3536  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  and HfB2 (MeB2 ) baseline UHTCs possess insufficient oxidation resistance for hypersonic flight environments, as they start oxidizing at temperatures as low as 800 C8 according to the following reaction:     MeB2(s) + 5/2 O2(g) → MeO2(s) + B2O3(s,l)  (1)                 Liquid B2O3 fills pores and protects the un-oxidized volume from further oxidation as it acts as a passive layer for temperatures above 1200 oxygen diffusing inwards. At C, the B2O3 is lost by evaporation and a columnar porous layer of MeO2 is left, resulting in continuous active oxidation of MeB2 .8 Therefore, hypersonic applications require materials that could overcome this limitation. Subsequently, much research has focused on development of more oxidation-resistant UHTCs composites.9-12 A promising approach has been the addition of SiC to MeB2 , typically between 10 and 30 vol.% SiC, in which a protective glassy borosilicate (BS) coating is formed in oxidizing environments.6,7 Pure SiO2 coatings melt at 1650 C. Although the BS coating melts at lower temperatures it could withstand temperatures up to 2000 C for up to 1 h.13-18 In addition, it should be noted that the BS layer is protective up to 2000 C in non-air flowing atmosphere.19 Nonetheless, at very high temperatures (>2000 C) the BS layer is lost and subsequently the underlying bulk material is exposed to continuous (active) oxidation. However, Han et al.20 showed that at high enough temperatures (2200 C) in an oxyacetylene torch test, the remaining porous ZrO2 layer could sinter forming a dense protective ZrO2 layer. However, oxygen diffusion through a dense ZrO2 layer is quicker than through a pure SiO2 layer. For example, the oxygen diffusion coefficient C is 10 −10 m2 /s21 while in pure SiO2 in ZrO2 at 1500 it is 10 −21 m2 /s at 1550 C.22 In addition, the activation energy for oxygen diffusion in ZrO2 is >200 kJ/mol23 and 120 kJ/mol for SiO2 24 which indicates that the difference in oxygen diffusion coefficient between ZrO2 and SiO2 increases with temperature. Therefore, the protective behavior of dense ZrO2 layer is questionable as it is highly permeable to oxygen. Oxidation of SiC-reinforced UHTCs is controlled by the BS coating as it is the oxidation limiting layer for oxygen diffusion. It is assumed that the BS coating is protective at least up to temperatures of 1650 C.25 However, features such as convection cells that represent a quick path for oxygen diffusion inwards to un-oxidized volume in addition to explosive bubble formation on the top surface at temperatures as low as 1550 C,22 indicate that the BS coating might not be purely protective, as has been assumed. Oxidation kinetics are usually analyzed by a power rate equation,26 as follows;                 (x − x0 )n = kt ↔ x = x0 + k (cid:7)  √  n  t  (2)  x  where is the mass change or oxide scale thickness, is the oxide layer thickness without holding time or a constant that accounts for oxidation effects far from equilibrium, k is a constant related to the oxygen diffusion coefficient and n is the exponent depending on the oxidation mechanism (n = 1, un-protective linear kinetics; n ≥ 2, protective kinetics as solid  x0  (cid:7)     three20,27,33,34 or  state diffusion is the limiting rate for oxidation). Most research to date on oxidation mechanisms in reinforced-UHTCs has focused on the study of oxide scale formed at different oxidation temperatures for a fixed time5,6,13-16,18,27-31 without any evaluation of the time evolution of the oxide scale. As a result these cannot be considered as oxidation kinetics studies. Other studies investigated the oxide scale for a fixed temperature and two different times.22,32 Some studies have investigated the oxidation kinetics evolution for a fixed temperature and four24 different times, giving a better indication of the oxide scales ability to protect. However, some of these studies are misleading as they include the datum point (0, 0) for weight gain and oxide scale at no holding time26,33 or they force fit curves to go through the origin, which is incorrect rates in open air box furnaces (10 as UHTCs oxidize during heating because of the slow heating C/min). Furthermore, Guo and Zhang26 included the datum (0, 0) and calculated kinetic exponents n of 2 for weight gain in ZrB2 -10 vol.% SiC and ZrB2 -30 vol.% SiC, 1 for BS glassy layer in ZrB2 -10 vol.% SiC, 2 for ZrB2 -30 vol.% SiC, 2 for extended SiC-depleted layer in ZrB2 -10 vol.% SiC and 3 for extended SiC-depleted layer in ZrB2 -30 vol.% SiC, which could be considered to confer a protective behavior. In fact, if the (0, 0) datum is not considered the calculated kinetic exponents n are 1 for weight gain in ZrB2 -10 vol.% SiC and ZrB2 -30 vol.% SiC, <1 for BS glassy layer in ZrB2 -10 vol.% SiC, 1 for ZrB2 -30 vol.% SiC, 1 < n < 2 for extended SiC-depleted layer in ZrB2 -10 vol.% SiC and 1 for extended SiC-depleted layer in ZrB2 -30 vol.% SiC, which is non-protective behavior.26 Moreover, there is experimental evidence that UHTCs oxidize without any holding time in open air furnace isothermal experiments.30,34 Other oxidation kinetics studies based on weight change analysis33,35 to support the argument that it provides protection, do not necessarily correspond to oxide scale results as a consequence of the volatilization of some species, such as B2O3 and CO2 . For example, Zhang et al.36 showed that weight gain kinetics have an oxidation exponent of 1 for pure ZrB2 and 2 for ZrB2 -4 wt.% WC while the oxidation exponents n for the oxide scale formation were <1 and 1 respectively. In addition, a dense oxide layer was reported as the porous ZrO2 layer was liquid sintered with WO glassy melts,37 which could not be beneficial for the system due to the high oxygen diffusion coefficient of ZrO2 resulting in a less oxidation resistant UHTC compared to an UHTC with a protective SiO2 layer on top. Different approaches have been tried to improve oxidation resistance of UHTCs containing SiC at extreme temperatures11 including (i) adding different Si-containing compounds with the aim of forming a solid solution with MeO2 and incorporating different elements in the BS coating, thus increasing its viscosity and melting temperature,26,33,35 (ii) adding different MeB2 phases such as TaB2 , TiB2 or CrB2 since borate and silicate glasses containing oxides of the elements listed (Group IV-VI transition metals) are immiscible and lead to phase separation, increasing viscosity and melting temperature,37 and (iii) formation of a protective refractory coating on UHTC oxidation as a consequence of the addition of some rare earth additives to SiC-reinforced ZrB2 after 1 h oxidation at 1600 C (Jayaseelan     \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3537           et al.28 ). In the particular case of La2O3 addition, the melting point of the crystal phase in the dense and solid protective coating formed (La2Zr2O7 ) is 2200 C, which makes it a promising candidate for UHTCs needing to withstand temperatures around 2000 C. There have been only a few studies of long-term oxidation of UHTCs; a couple below 1400 C24,29 and one for up to 10 h at 1500 C,27 which motivated this study. The goal of this work is to study the long-term oxidation kinetics of UHTCs at 1400 C, 1500 C and 1600 C for oxidation times of up to 32 h, 16 h and 4 h, respectively in order to (i) determine whether the BS coating is protective on SiC-reinforced ZrB2 and HfB2UHTCs and (ii) study oxidation kinetics of La2O3 -doped SiC-reinforced ZrB2 and HfB2 UHTCs and analyze the protectiveness of the oxide products formed during oxidation.              2. Experimental procedure  d50  2.5 \\u242em, ZrB2 powder (>99%, ρ = 6.085 g/cm3 , Sigma Aldrich, Gillingham, UK), HfB2 powder (>99%, d50  5.0 \\u242em, ρ = 10.5 g/cm3 , Sigma Aldrich, Gillingham, UK), SiC powder (␣-SiC, 99%, d50  0.7 \\u242em, ρ = 3.217 g/cm3 , Good Fellow Chemicals, Huntingdon, UK) and La2O3 d50  10 \\u242em, (>99%, ρ = 6.51 g/cm3 , Fluka chemicals supplied through Sigma Aldrich, Steinheim, Germany) were used to form different UHTC compositions; ZrB2 + 20 vol.% SiC (ZS20), ZrB2 + 20 vol.% SiC + 2 wt.% La2O3 (ZS20La), HfB2 + 20 vol.% SiC (HS20) and HfB2 + 20 vol.% SiC + 2 wt.% La2O3 (HS20La). Ceramic powders were processed and then sintered using Spark Plasma Sintering (SPS) as described previously.10 The main microstructural features of SPS UHTCs, as shown in previous work10 and Fig. 1, are; (I) high density (>99% of theoretical relative density); (II) ZrB2 -based UHTCs contain a homogeneous dispersion of SiC throughout the ZrB2 matrix with SiC grain sizes between 1 and 2 \\u242em while HfB2 -based UHTCs contain an inhomogeneous dispersion of SiC throughout the HfB2 matrix with SiC agglomerates as large as 20 \\u242em; (III) La2O3 particles are often in close proximity to SiC particles. Large agglomerates of SiC arise from the density difference between SiC (ρ = 3.217 g/cm3 ) and HfB2 (ρ = 10.5 g/cm3 ), which promotes their non-homogeneous mixing during thermo-rotary drying of HfB2 -based UHTCs solutions. The use of a surfactant might overcome this limitation and is being investigated. HS20La contains some SiC agglomerates of 10 \\u242em, which is about half the size of the HS20 agglomerates. The location of La2O3 (ρ = 6.51 g/cm3 ) particles near to SiC particles and its closer density to SiC than HfB2 assist particle dispersion in the solutions reducing the formation of SiC agglomerates. However, SiC distribution is more homogeneous throughout ZrB2 matrix than HfB2 matrix due to lower density difference between SiC (ρ = 3.217 g/cm3 ) and ZrB2 (ρ = 6.085 g/cm3 ) than between SiC (ρ = 3.217 g/cm3 ) and HfB2 (ρ = 10.5 g/cm3 ), which prevents second phase particles agglomeration in UHTCs solutions. Samples for oxidation tests were cut by electrical discharge machining (EDM) with the following dimensions; 25 mm × 2 mm × 1.5 mm for length, height and thickness. Oxidation tests for all compositions were carried out in an open air              furnace operating in isothermal mode at 1400 C for 0 h, 1 h, 4 h, 16 h and 32 h, 1500 C for 0 h, 1 h, 4 h, 8 h and 16 h and 1600 C for 0 h, 1 h, 2 h, 3 h and 4 h. Heating and cooling rates were 10 K/min and 20 K/min respectively. Longer times at 1600 C were not studied to avoid catastrophic failure of the Superkanthal heating elements. Experiments under the same conditions were made once for all compositions so that all data are directly comparable. Samples were placed on a high purity Al2O3 cruof 0.5 mm separated by 20 mm, minimizing the contact area, cible (>99.99%) in contact with their two edges with a thickness possible reactions and glass flow between the UHTCs and the crucible. A platinum foil was placed between the samples and the crucible to avoid chemical reactions between the crucible and the samples. Mass gain after oxidation was calculated by weighing before and after oxidation. Weighing was carried out by; (i) weighing of samples with platinum foils and crucible, before and after oxidation; (ii) weighing of samples before and after oxidation to ensure mass gain measurements are just related to oxidation of UHTCs instead of reactions between samples and crucible and quantify possible glass flow exchange between samples, platinum foils and the crucible. Cross section specimens for SEM were prepared using conventional methods involving successive steps of grinding and polishing with diamond slurries and cloths embedded with up to 1 \\u242em diameter particles. The specimens were observed in an SEM equipped with a field emission gun (LEO 15, JEOL, Tokyo, Japan). Samples were examined in secondary electron (SE) imaging mode for oxide scale thickness quantification and backscattered electron (BS) image mode for phase morphology determination. Energy dispersive X-ray spectroscopy (EDS) was used for phase identification.  3. Results and discussion  3.1. Oxidation kinetics characterization                 Oxidation kinetics at 1400 C, 1500 C and 1600 C are illustrated in Fig. 2, in which mass gain per unit area and oxide scale thickness versus time are represented. In addition, best fitting curves according to the power law equation (2) are shown. Oxidation kinetics exponents n are summarized in Table 1, which indicates a mix of pure parabolic behavior (n  2) for almost all compositions and linear and parabolic behavior (1 < n < 2) or conditions, except for HS20 at 1500 C and 1600 C (highly unprotective or unstable behavior n  0.5), ZS20La at 1600 C (n  8, highly protective behavior) and HS20La at 1400 C (n  4, highly protective behavior). Moreover, oxidation kinetic exponents n in mass gain were found to be similar to those determined from oxide scale thicknesses, except for ZS20La at C (n  2 and n  8, respectively) and HS20La at 1600 1600 C (n  1 and n  2, respectively), which may indicate mass adsorption in the oxide layer as mass gain increases faster than oxide scale thickness. Fig. 3 shows the oxidation of all UHCTs compositions at 1400 C without holding time to illustrate the initial oxidation of UHTCs, justifying use of the term x0 in fitting curves. Therefore, the inclusion of point (0, 0) in previous oxidation kinetics studies is misleading as oxidation takes place during heating as                \\x0c', '3538  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  Fig. 1. SEM pictures of UHTCs: (a) ZS20, (b) HS20, (c) ZS20La and (d) HS20La. Light gray particles next to SiC are La2O3  Reprinted from Ref. 10 with permission 2013, Journal of the European Ceramic Society.  a consequence of low heating rates (10 K/min) in open air box furnace experiments without any holding time.30,34 This effect is also far from equilibrium and addressed by the term x0 .  3.2. Oxidation kinetics review  activation energy for oxidation (120 kJ/mol)24 carbide has been found to be equal to the silicon activation energy,39 which suggests the same oxidation mechanism is present in both oxidation processes. Silicon carbide oxidizes according to the following reaction,  The oxidation rate of silicon by water or oxygen was first modeled by Deal and Grove38 via the following equation,  x2 + Ax = B(t + τ )  (3)  t  where x is the oxide scale thickness, is the time of oxidation and A, B and τ are coefficients independent of x and t but which can be functions of gas pressure, temperature and oxygen diffusion coefficient. Eq. (3) can be derived by assuming there are two processes involved in the oxidation, i.e. diffusion through the oxide film and a chemical reaction at the silicont (cid:10) a, where a is oxide interface. If the time for SiO2 layer growth just by chemical reaction at the interface without a significant diffusion contribution, linear behavior is expected while a parabolic behavior is expected for t (cid:11) a. Short times refer to the initial stage of oxidation where oxidation scale thickness is on the nanoscale. Therefore parabolic behavior is expected at the macroscale (oxidation scale thickness >1 \\u242em). Silicon  SiC(s) + O2(g) → SiO2(s) + C(s) C(s) + O2(g) → CO2(g)  (4)  (5)     It is usually assumed that oxidation of solid carbon is instantaneous and immediately vaporizes as CO2 . However, Ni et al.30 studied the initial stage of sintering at 1500 C, observing in the interface between un-oxidized HS20 volume and the SiCdepleted region formation of C particles encapsulating SiC. Therefore, it is more appropriate to separate Eqs. (4) and (5). Only passive oxidation of SiC is considered in this study as in the temperature range studied here, active oxidation of SiC, which involves complete SiO2 volatilization, is not expected and does not take place. The Deal and Grove model has been improved by considering the effect of a reaction layer,40 which adds a logarithmic term to Eq. (3), and ionic growth models for oxide scale with electric field-dependent ionic conduction,41 which leads to power law equations for Eq. (2) with oxide kinetic exponent n  Table 1  Oxidation kinetics exponent for mass gain and oxide layer thickness at 1400     C, 1500     C and 1600     C for ZS20, ZS20La, HS20 and HS20La.  Material  ZS20  ZS20La  HS20  HS20La  n (1400     C)  Mass gain 1.2 ± 0.1 1.8 ± 0.3 2.3 ± 0.3 3.2 ± 0.6  Oxide layer 1.4 ± 0.3 1.8 ± 0.3 2.2 ± 0.4 3.7 ± 0.8  n (1500     C)  Mass gain 2.0 ± 0.4 2.4 ± 0.3 0.4 ± 0.1 1.9 ± 0.2  Oxide layer 1.6 ± 0.1 2.6 ± 0.5 0.5 ± 0.1 2.2 ± 0.3  n (1600     C)  Mass gain 1.4 ± 0.1 2.1 ± 0.2 0.6 ± 0.1 1.4 ± 0.1  Oxide layer 1.5 ± 0.1 7.5 ± 0.4 0.5 ± 0.1 2.2 ± 0.2  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3539  0  4  8  12  16  20  24  28  32  0  5  10  15  20  25  30  35  M  a  s s  a g  i  n  /  n u  i  t  a  r  a e  (  m  g  /  c  m  ) 2  Time (h)   ZS20   ZS20La   HS20   HS20La  0  4  8  12  16  20  24  28  32  0  100  200  300  400  500  600   ZS20   ZS20La   HS20   HS20La  O  x  i  e d  l  a  y  e  r  t  h  i  k c  e n  s s  (  µ  m  )  Time (h)  0  2  4  6  8  10  12  14  16  0  2  4  6  8  10  12  14  16  M  a  s s  a g  i  n  /  n u  i  t  a  r  a e  (  m  g  /  c  m  ) 2  Time (h)   ZS20   ZS20La   HS20   HS20La  0  2  4  6  8  10  12  14  16  40  80  120  160  200   ZS20   ZS20La   HS20   HS20La  O  x  i  e d  l  a  y  e  r  t  h  i  k c  e n  s s  (  µ  m  )  Time (h)  0  1  2  3  4  0  2  4  6  8  10  12  14  16  18  20  M  a  s s  a g  i  n  /  n u  i  t  a  r  a e  (  m  g  /  c  m  ) 2  Time (h)   ZS20   ZS20La   HS20   HS20La  0  1  2  3  4  0  40  80  120  160  200  240  280  320  360   ZS20   ZS20La   HS20   HS20La  O  x  i  e d  l  a  y  e  r  t  h  i  k c  e n  s s  (  µ  m  )  Time (h)  A  B  C  D  E  F  1400 ºC  1400 ºC  1500 ºC  1500 ºC  1600 ºC  1600 ºC  Fig. 2. Mass gain and oxide scale layer thickness versus time at: (a) and (b) 1400     C, (c) and (d) 1500     C and (e) and (f) 1600     C, respectively, for all UHTCs  compositions. Best fitting to Eq. (2) is drawn.  from 1 to 4 although in the case of SiO2 formation, an oxide kinetic exponent from 1 to 2 is expected.42 Other features that can alter oxidation kinetics are the formation of an open porous lower the oxidation kinetics exponent (n ≤ 1).43,44 For example, layer, cracks or spallation, which enhance oxidation rate and Zhang et al.36 showed oxidation kinetic exponent <1 for oxide scale thickness of pure ZrB2 oxidized at 1500 C. SiC-reinforced     UHTCs oxidation implies more phenomena are involved than in oxidation of pure silicon or SiC as there is a porous oxide layer under the BS layer and a SiC-depleted layer. However, the rate limiting diffusion process is the diffusion through the BS layer as diffusion through a porous ZrO2 layer is quicker. In fact, the oxyC is 10 gen diffusion coefficient of SiO2 at 1550 is 10 −10 m2 /s.21 Therefore, while for ZrO2 it the SiO2 layer     −21 m2 /s22                                      \\x0c', '3540  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  Fig. 3. SEM pictures of initial stage of oxidation at 1400     C without holding time, which justify the use of term x0 in Eq. (2), for: (a) ZS20, (b) ZS20La, (c) HS20  and (d) HS20La.  is more protective than the ZrO2 layer and controls oxidation as it is the oxygen diffusion limiting path. Generally, an oxidation kinetics exponent between 1 and 2 is expected in SiC-reinforced UHTCs in an intermediate temperature range as studied here since the BS coating is a melt and fully solid state diffusion is not the limiting rate, in agreement to results shown in Table 1 in which oxide layer oxidation kinetic exponent is 1 ≤ n ≤ 2 for SiC-reinforced UHTCs (ZS20 and HS20).  3.3. Oxidation kinetics discussion  3.3.1. ZS20  Oxidation of ZS20 is explained in terms of an oxidation power law with exponent n from 1 to 2 in the temperature range 1400-1600 C. In addition, mass gain and oxide scale thickness of the BS layer is larger at 1400 C than at 1500 C for the same oxidation time as seen in Figs. 2 and 4 respectively, which is explained in terms of lower boria content, which vaporizes from the BS layer more rapidly at 1500 C than at 1400 C. In fact, oxygen penetrates deeper in ZS20 at 1400 C than at 1500 C as the oxygen diffusion coefficient is greatly reduced by boria evaporation, as seen in Fig. 5 which reveals a much thicker scale at 1400 than 1500 C. According to Opeka et al.3 , B2O3 content in the BS melt decreases with temperature and approximately stabilizes in the range 1200-1400 to a 90 wt.% (89 mol%) SiO2 BS melt. C to a level of 10 wt.% B2O3 , which corresponds However, there are no data on the boria content after oxidation at 1500 C when presumably B2O3 content in the BS melt should have been reduced. Considering a SiO2 concentration increase of 5 mol% as a consequence of B2O3 evaporation, which is reasonable, the oxygen diffusion coefficient through the BS melt is reduced 5 times at 1550 C.22 Moreover, considering an activation energy of 125 kJ/mol for oxidation, it is calculated that = 2DO2 1400  C . Therefore, a layer of BS melt with                                   DO2 1500  C  0  4  8  12  16  20  24  28  32  0  100  200  300  400  500  600   Oxide scale thickness   BS layer thickness  SiC-depleted + porous ZrO  2   thickness  O  x  i  e d  c s  a  l  e  t  h  i  k c  e n  s s  (  μ  m  )  Time (h)  0  2  4  6  8  10  12  14  16  0  100  200  300  400  500  600   Oxide scale thickness   BS layer thickness  SiC-depleted + porous ZrO  2   thickness  O  x  i  e d  c s  a  l  e  t  h  i  k c  e n  s s  (  μ  m  )  Time (h)  A  B  1400 ºC  1500 ºC  Fig. 4. Oxide scale thickness of ZS20 from different  layers (total oxide scale,  BS melt and porous ZrO2 + SiC-depleted layers) at: (a) 1400     C and (b) 1500     C.              \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3541  Fig. 5. SEM micrographs of oxide scale of ZS20 for a holding time of 16 h at: (a) 1400     C and (b) 1500     C (c) magnified image of (a) for direct eye comparison with  (b).                                      94 mol% SiO2 at 1500 C could be as protective as an 89 mol% SiO2 layer 2.5 times thicker at 1400 C. For example, a 100 \\u242em thick BS melt layer containing 89 mol% SiO2 at 1400 C would be as protective as a 100 \\u242em thick BS melt layer containing C (2.8 mol% SiO2 increase). Fig. 4 91.8 mol% SiO2 at 1500 shows the thickness of BS melt and porous oxide layer at 1400 C (Fig. 4a) and 1500 C (Fig. 4b) for ZS20, which illustrates that only a B2O3 content reduction could explain differences in oxidation behavior. The observed higher curvature of the BS melt surface at 1400 C (Fig. 5a) than at 1500 C (Fig. 5b), indicates that viscosity increases with temperature, consistent with a reduction of B2O3 in the BS melt. Moreover, BS melt thickC is 2.5 times thicker ness at 1400 than BS melt thickness at 1500 C as observed in Fig. 5, which suggests that B2O3 volatilization could be increased around 5 mol% from 1400 C to 1500 C. Moreover, at 1600 C the vaporization of boria from the BS layer is sufficient to produce a decrease of BS scale thickness with time as shown in Fig. 6, which indicates that BS vaporization is more rapid than BS formation. In addition, the BS melt mixes more homogeneously with the porous oxide layer at 1600 C as a consequence of increasing BS melt viscosity due to further B2O3 reduction. Low viscosity of BS layer causes the oxidation response to deviate from the purely parabolic behavior of Deal and Grove38 to a mix of linear-parabolic terms. In fact, bubbles in addition to convective cells, which could be considered as system instabilities, were identified by SEM at temperatures as low as 1500 C, which is in close agreement with previous observations of ZS15 oxidation by Karlsdottir and Halloran at 1550 C22 and is clearly associated with boria evaporation from the BS melt.              3.3.2. ZS20La  C (n  2) ZS20La showed pure parabolic behavior at 1400 and its oxidation kinetic exponent increases with temperature to                             2.5 at 1500 C. In addition, values of n for mass gain and oxide scale thickness at 1600 C are different by almost four times, 2.1 for mass gain and 7.5 for oxide scale, which indicates that an oxygen adsorption mechanism may be active because large changes in mass gain lead to small changes in oxide scale. Oxide scale thickness of ZS20La is smaller than oxide scale thickness of ZS20 at 1400 C while the converse applies for 1500 C and 1600 C as seen in Fig. 2, presumably due to the higher volatilization of B2O3 present in the BS layer of ZS20 at these temperatures. Therefore, ZS20La is more oxidation resistant at 1400 C but not at 1500 C and 1600 C. The main difference between the oxide layers of ZS20 and ZS20La is the mixing of the BS glassy melt with the ZrO2 scale which is more viscous in ZS20La and does not ascend directly to the top surface, producing a more homogeneous mixing between glassy and oxide phase as seen if Fig. 7 is compared to the ZS20 oxide layer in Fig. 5a or b. Homogeneous mixing of BS melt with ZrO2 could be detrimental for the system as diffusion through ZrO2 is much quicker than through a BS melt even if it has higher viscosity due to the presence of La2O3 . La2O3 proximity to SiC promotes the inclusion of La2O3 into the BS phase, increasing viscosity and thermal stability as well as lowering the oxygen diffusion coefficient through the glassy phase. In fact, La2O3 is used as an additive to increase creep resistance of MoSi2 ,45 which contains glassy SiO2 at grain boundaries and triple junctions. Moreover, the enhancement of creep resistance was related to the following reaction;  2SiO2(s) + La2O3(s) → La2Si2O7(g)  (6)  Reaction was complete or SiO2 was completely eliminated at a La2O3 /SiO2 wt.% ratio of 3.2.45 Two different polymorphs of La2Si2O7 were observed with a wt.% ratio of 6.4, which corresponds to a 4 wt.% La2O3 /MoSi2 composites (5 vol.%). Therefore, in the present study La2O3 is completely dissolved in  \\x0c', '3542  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  Fig. 6. SEM micrographs of oxide scale of ZS20 at 1600     C for a holding time of: (a) 1 h and (b) 2 h.  Fig. 7. SEM micrographs of oxide scale of ZS20La at 1500     C for a holding time of: (a) 8 h and (b) 16 h. Arrow indicates precipitation of La2O3 -SiO2 particles.  the La2O3 -doped UHTCs BS melt for <5 vol.% La2O3 content and protective Me2La2O7 is not formed unless La2O3 concentration is at least 10 vol.%.28 However, precipitates of La2O3 -SiO2 phase were found on top of the oxide layer and identified with EDS in samples oxidized at 1500 C as observed in Fig. 7a indicated by the arrow, which confirms higher B2O3 volatilization at 1500 C than at 1400 C, facilitating the precipitation of La2O3-SiO2 particles. Furthermore, a completely new phenomena was observed on samples oxidized at 1600 C.              ZrOxCy particles were detected by EDS within the BS melt and on top of the oxide layer, as shown in Fig. 8a and b. This top ZrOxCy gets oxidized with time and formed ZrO2 , while intermediate ZrOxCy still remains deeper in the oxidized layer after long exposures times. Similar behavior was found for formation of SiOxCy particles. Therefore, the ZrOxCy and SiOxCy particles formed act as an active-protective barrier that becomes further oxidized after long exposures times, which could account for the high oxidation protective behavior observed in oxidation  Fig. 8. SEM micrographs of oxide scale of ZS20La at 1600  C for a holding time of: (a) 0 h and (b) 2 h. EDS spectra of (c) ZrOxCy and (d) SiOxCy . Some ZrO2 grains are rounded in black circles and ZrOxCy are rounded in white circles.     \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3543  kinetics for ZS20La. The ZrO2 particles could react with C(s) after SiC oxidation produced according to Eq. (4).  ZrO2(s) + xC(s) → ZrO2Cx(s)  (7)  Detection of C by EDS in the SEM is difficult and phase identification is uncertain. Different compositions were quantified by EDS, such as ZrO2C3 , Zr2O3C4 , ZrO5C4 and ZrO4C6 . The only trend observed is that after long exposure times the very top surface comprises ZrO2 particles, whereas C concentration in ZrO2Cx does not follow any trend and C atomic concentration ranges from 40 to 60 atomic %. The mechanism according to reaction (7) is that ZrO2Cx particles are formed on the oxide layer interface and become oxidized on the top surface after long exposure times. Therefore, C concentration on ZrO2Cx particles should increase from top layer to oxide layer interface. However, this trend is not observed and there is not a clear pattern about C concentration throughout the oxide layer, which may indicate there is another active mechanism to form ZrOxCy particles. Another possible route for the formation of ZrOxCy could be described as follows; There are two species containing O2(g) being transported through the BS layer, O2(g) from the atmosphere and CO2(g) released from SiC(s) or ZrOxCy active oxidation, according to Eqs. (4), (5) and (7). As there is no report of C(s) in BS layer, the following mechanism is suggested as a possible reaction to form ZrOxCy from dissolved Zr in BS melt and released CO2(g) ;  Zr (l) + xCO2(g) → ZrO2xCx(s)  (8)     The presence of La2O3 within the BS melt increases its viscosity and reduces the O2 and CO2 diffusion coefficients, so promoting reaction (8) to take place at 1600 C. In addition, ZrO2xCx could react and be oxidized by atmospheric O2 forming an active-protective barrier for inwards oxygen diffusion. The adsorption of species is also noticeable and confirmed by the difference between mass gain kinetics and oxide layer thickness kinetics. Moreover, a large mass gain change produces a relatively small increase of oxide scale thickness as a consequence of oxidation kinetics exponent between mass gain and oxide layer (2 and 8, respectively), as seen in Fig. 2. An oxidation kinetic exponent as high as 8 could be obtained if Fick’s law for atomic diffusion is constrained by the presence of a particle sink.46 Furthermore, the more particles formed within the oxide scale, the more limited the oxygen diffusion. The continuous oxidation of ZrOxCy and SiOxCy particles and continuous recombination or possible adsorption of oxygen-containing species explain how an appreciable change in mass gain produce a small change in oxide layer thickness, addressing the difference in measured values for oxidation kinetics exponents (Table 1). However, the spatial distribution of ZrO2 and ZrOxCy is not clear as it is likely there is more than one mechanism forming ZrOxCy ; this is the subject of further investigation. Therefore, ZrOxCy coatings are presented as new candidates for UHTCs protection against oxidation. Fig. 9 shows an overview of ZS20La oxide layer formed at 1600 C after 4 h oxidation. Precipitation of ZrSiO4 was detected and characterized by small light contrast features at micron scale level (0.5 to 2 \\u242em)     Fig. 9. SEM micrograph of oxide scale of ZS20La at 1600  C for 4 h. Some ZrO2 grains are rounded in black circles and ZrOxCy are rounded in white circles.     in the BS layer and on the outer surface of ZrB2 -based UHTCs as well as drag of ZrO2 particles into the BS layer. These particles are observed in Figs. 5-9 being more evident for the conditions with longer oxidation dwell. Crystallization of SiO2 was not detected even for long exposures times as it is a melt during oxidation and the glassy character of the BS layer in ZrB2 -based ceramics is clearly visible in microstructures of oxide layer. In addition, crystallization phenomena usually occur under equilibrium conditions which are not present in this study due to continuous precipitation, vaporization or diffusion of different species during oxidation.  3.3.3. HS20     HfB2 has a lower reactivity with oxygen and HfO2 has lower oxygen permeability compared to ZrB2 and ZrO2 , respectively.9 As a result B2O3 content in the BS melt in HS20 samples may be lower or even negligible compared to ZS20. Therefore, a much thinner BS layer protects the HfB2 -based UHTCs more effectively than in ZrB2 -based UHTCs, as seen in Fig. 2. HS20 at 1400 C (Fig. 2) show a pure parabolic-protective behavior as a consequence of the lower B2O3 content in its BS melt. However, there is an important microstructural difference between ZS20 and HS20, which is that the distribution of SiC particles throughout the HfB2 matrix is not homogeneous and there are large SiC agglomerates (up to 20 \\u242em dia.). Therefore, once the oxide scale is thicker than 20 \\u242em some SiC agglomerates have been already oxidized leaving larger cavities, up to five times larger along one direction and up to 20 times larger in terms of surface covered by SiC agglomerate as seen in Fig. 1. Their presence has the following effects; (i) capillary forces that move BS melt to the top are lower and then oxidation protection is less efficient, and (ii) the generation of a large cavity within a porous layer could be detrimental as it communicates with other pores and could lead to detachment of oxide particles or enhance oxidation system instabilities. As a consequence, the mass gain and oxide layer kinetic exponents are 0.5 at 1500 C and 1600 C. This does not mean that the BS coating is completely unprotective. It indicates that for short periods of time (<4 h at 1500 C and <2 h at 1600 C) the behavior is protective and after a critical oxide layer              \\x0c', '3544  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  Fig. 10.  (a) SEM micrograph of oxide scale of HS20La at 1600     C for 3 h illustrating the presence of different phases of HfOxCy and SiOxCy detected by EDS in  (b) near the top exposed surface and (c) interface with un-oxidized UHTC volume.  thickness is reached, the oxidation behavior turns into an unprotective linear behavior, which as a result gives a mean kinetic exponent of 0.5 for the whole process, as depicted in Fig. 2. Therefore it is important to homogenize the SiC distribution throughout UHTCs matrix during processing to avoid unstable oxidation behavior.  area near the top surface layer and oxide layer interface with unoxidized volume are magnified in Fig. 10b and c, respectively. Oxidation mechanisms are slightly different than in ZS20La because in HS20La SiC gradually oxidized and there are small  3.3.4. HS20La        HS20La behaves similar to ZS20La, with the difference that it is not more oxidation resistant than HS20 at 1400 C due to lower content of B2O3 in the BS layer in HfB2 -based UHTCs. However, the presence of SiC agglomerates did not affect the stability of its behavior at 1500 C and 1600 C, which is protective as shown in Table 1 and Fig. 2. In addition, the oxide layer thickness kinetic exponent at 1600 C is smaller than ZS20La, which could be related to either the presence of SiC agglomerates, the lower oxidation rates of HfB2 -based UHTCs compared to ZrB2 -based UHTCs or a lack of evidence of BS melt phase in the oxide layer. Furthermore, the formation of different HfOxCy phases was detected by EDS. Fig. 10a shows the oxide scale layer of HS20La after oxidation at 1600 C for 3 h, in which an           Fig. 11. BS SEM micrograph of HS20La at 1600  with un-oxidized UHTC volume.     C for 3 h near the interface  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3545  0  1  2  3  4  5  6  7  8  9  10  Si  3  O  2  C  4  CO  Si  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  0  1  2  3  4  5  6  7  8  9  10  Hf  2  O  6  C  3  Hf  Hf  Hf  Hf  O C  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  0  1  2  3  4  5  6  7  8  9  10  HfO  2  O  Hf  Hf  Hf  Hf  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  0  1  2  3  4  5  6  7  8  9  10  HfO  3  Hf  Hf  Hf  O  Hf  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  0  1  2  3  4  5  6  7  8  9  10  HfOC  Hf  Hf  Hf  O C  Hf  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  0  1  2  3  4  5  6  7  8  9  10  HfO  4  C  24  C  O  Hf  Hf  Hf  C  n u o  t  s  (  A  r  b  .  n u  i  t  s  )  Energy (keV)  A  B  C  D  E  F  Fig. 12. EDS spectra of different phases detected after oxidation at 1600     C for 3 h in HS20La.  SiC particles (0.5 2 \\u242em) clearly visible at the interface (Fig. 10b). In addition, a new oxide phase, HfO3 , was detected. A BS image (Fig. 11), reveals different contrast between HfO2 and HfO3 as well as with La-SiOxCy particles, small particles of HfOxCy (5 10 \\u242em) with high (85 atomic %) C content and small (0.5 2 \\u242em) SiC particles in the oxide layer. Different EDS  spectra from the phases in HS20La are illustrated in Fig. 12. The C content of HfOxCy particles range from 30 to 85 atomic %, which is a broader range than the observed in ZS20La that might be related with a higher reactivity of HfO2 with C or thermal stability in air of HfOxCy particles. Another interesting feature is the absence of a clear BS phase which led to the formation of                          \\x0c', '3546  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  Fig. 13.  (a) BS SEM micrograph of HS20La initial stage of oxidation at 1600     C without holding time with (b), (c), (d) EDS spectra of different particles formed  during oxidation.           more oxidation products than in ZS20. Ni et al.30 detected C particles at the interface between oxide scale and un-oxidized volume during the initial stage of oxidation at 1500 C. The issue is why this phenomena was not observed in ZS20La. This could be explained because the BS melt of ZrB2 based ceramics has higher B2O3 content and there is no clear evidence of BS melt in HS20La above 1500 C, which lowers oxygen diffusion coefficient. In addition, 2 wt.% of La2O3 addition is equivalent to a higher volume content of La2O3 in the BS melt of HfB2 -based UHTCs than in ZrB2 -based UHTCs, which lowers the oxygen diffusion coefficient. As a consequence, C volatilization kinetics at the interface are lowered as oxygen diffusion is largely reduced, which promotes C reaction with other surrounding species. Therefore, a low addition of La2O3 leads to the formation of different oxycarbide compounds of Si, La-Si and Hf at 1600 C throughout the oxide scale layer, which might continuously react and get oxidized with O2 slowing the diffusion of oxygen-containing species and promoting this new protective mechanism against oxidation. The initial stage of oxidation was investigated in HS20La at 1600 C (Fig. 13). HfC was detected on the top exposed surface and below the HfC a thin (aprox. 5 \\u242em) layer of HfO3 was detected on top of HfO2 . HfO2 adsorbing O2 to form HfO3 is an unlike mechanism. However, a chemical reaction with available C is more likely which may lead to the formation of the detected particles as follows; 3HfO2(s) + C(s) → HfC(s) + 2HfO3(s)  (9)     In addition, formation of La2O3 -SiO2 and La-SiOxCy particles was detected. Furthermore, in the un-oxidized volume La2O3 was seen to react with SiC forming a La-SiOxCy  particle with higher C content (Fig. 13). Whether the formation of these La-SiOxCy particles is beneficial or not for UHTCs bulk properties is under investigation. An ongoing TEM study is being carried out to determine exact stoichiometry, crystalline nature and localization of different phases formed as well as a more precise chemical analysis to determine the C content of different particles formed with higher accuracy. La2O3 was added to MeB2-SiC composites with the intention of forming protective Me2La2O7 . However, La2O3 dissolved in the BS melt at low La2O3 contents (<5 vol.%), which, however, promoted new oxidation reactions with protective behavior. As a result, MeOxCy phases are suggested as new protective coatings for UHTCs as oxygen diffusion coefficients are reduced and oxide scale evolution is retarded with a stabilization trend for long exposures times, which is a direct consequence of the formation of MeOxCy particles. The more MeOxCy particles formed, the more oxidation resistant the UHTCs composites. The observed trend of La2O3 -doped UHTCs over long oxidation exposures times is that they are more oxidation resistant than SiC-reinforced UHTCs. MeOxCy are preferred for UHTC coating compared to SiOxCy as MeOxCy are more compatible with the UHTCs matrix and also might possess higher thermal stability.  4. Conclusions  Oxidation kinetics have been studied for ZrB2 and HfB2 based UHTCs. A partially protective behavior against oxidation was characterized for SiC-reinforced UHTCs (1 < n < 2). Moreover, the presence of SiC agglomerates led to a completely  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 34 (2014) 3535-3548  3547  unstable and unprotective behavior for long exposures times at 1500 C and 1600 C in HS20. La2O3 addition is beneficial for oxidation protection as it increases BS melt viscosity and lowers the oxygen diffusion coefficient through the BS melt. On the other hand, an increase of BS melt viscosity led to a more homogeneous mixing between oxide products and BS melt as well as reducing the thickness of the outer BS protective coating, which is detrimental for UHTCs oxidation resistance as it is the oxygen diffusion limiting layer. However, new reactions were observed at 1600 C to form MeOxCy with a more stable and protective behavior for long exposure times than SiC-reinforced UHTCs. Therefore, it is proposed that MeOxCy -containing coatings are promising and protective coatings for UHTCs, even after longterm exposures. The detection of different phases suggests that the O/C ratio could be tailored to minimize thermal mismatch damage at the interface arising from e.g. thermal fatigue.           Acknowledgments  The Authors’ acknowledge Prof. Mike Reece, Nanoforce Technology Ltd., Queen Mary, University of London, UK for providing access to the Spark Plasma Sintering facility. 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},{
  "_id": 45,
  "PDF": "Effect of LaB6 additions on densification, microstructure, and creep with oxide scale formation in ZrB2-SiC composites sintered by spark plasma sintering.pdf",
  "Text": "['Journal of the European Ceramic Society 39 (2019) 2782-2793  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e lsev ie r .com / loca te / jeu rce ramso c  Original Article  Effect of LaB6 additions on densification, microstructure, and creep with oxide scale formation in ZrB2-SiC composites sintered by spark plasma sintering  Sunil Kumar Kashyap, Rahul Mitra⁎  Department of Metallurgical and Materials Engineering,  Indian Institute of Technology Kharagpur, Kharagpur, 721302, West Bengal,  India  T  A R T I C L E  I N F O  Keywords: Zirconium diboride Spark plasma sintering Microstructure Creep Oxidation  A B S T R A C T  A comparative study has been carried out on densification, microstructure, and creep with oxide-scale formation in ZrB2-20 vol.% SiC-(7, 10 or 14 vol.%) LaB6 composite containing B4C and C as additives, and prepared by spark plasma sintering at 1800 °C under 70 MPa ram pressure. Addition of LaB6 has promoted densification of composites by scavenging oxygen impurity, thereby increasing their hardness. Constant load compressive creep tests at 1300 °C under 47 and 78 MPa stresses have shown the lowest creep rate in the 10 vol.% LaB6 composite. The stress exponents obtained for composites having 10 vol.% LaB6 (1.3 ± 0.1) and 14 vol.% LaB6 (2.6 ± 0.2) suggest respectively, grain boundary diffusion with intergranular glassy phase formation and dislocation glide as operating mechanisms. Intergranular cracking caused by grain boundary sliding appears as the damage mechanism. Oxide scales formed during creep exhibit greater thickness and defect concentration than those by isothermal exposure at 1300 °C within similar duration.  1.  Introduction  Ceramics based on the transition metal borides, nitrides, and carbides have extremely high melting points (> 2500 °C) and are referred to as ultra-high temperature ceramics (UHTCs). In recent years, the development of ZrB2 based UHTCs has received significant attention worldwide due to their unique properties including high melting temperatures, impressive creep and oxidation resistance, desirable thermal shock resistance due to the combination of low thermal expansion coefficient and high thermal conductivity, as well as their ability to withstand extreme environmental conditions [1]. Due to the combination of the properties mentioned above, the UHTCs are considered as preferred candidate materials for applications in leading edges and nose-cones of hypersonic flight and atmospheric re-entry vehicles, as well as rocket nozzles [2,3]. The conventional structural materials under load usually succumb in this kind of extreme environment to damage by oxidation and creep, which are predominant on exposure at high temperatures. The addition of SiC as reinforcement to ZrB2 matrix enhances the thermal conductivity, oxidation, and ablation resistance, whereas its theoretical density and coefficients of thermal expansion are reduced [4-6]. Presence of SiC as reinforcement helps in the formation of a protective borosilicate scale, which contributes to its oxidation resistance on exposure at elevated temperatures. Furthermore,  earlier studies have shown that grain-coarsening of the ZrB2 matrix is arrested through pinning by the SiC particles present as dispersoids [7,8]. Typically, higher hardness and flexural strength are accomplished with decrease in ZrB2 and SiC grain sizes, but this is achieved at the expense of the high-temperature creep resistance [9,10]. An earlier study on pressureless sintered ZrB2-SiC composites has shown the optimum volume fraction of SiC as 20% for providing the best combination of resistance to thermal shock and ablation [11]. One of the problems usually encountered regarding the development of ZrB2-SiC composite is its poor sinterability due to strong covalent bond and low diffusivity of both these phases. The addition of B4C and C as additives has shown a significant increase in relative density to > 99% at 1900 °C or 2000 °C by removing the surface oxides, i.e., ZrO2 and B2O3 from the powder particles, and hindering grain growth [12]. The addition of B4C not only enhances the densification, but also increases hardness, flexural strength, and wear resistance of the sintered composites [13,14]. Moreover, the presence of B4C may also contribute to the formation of B2O3 as a constituent of the oxide scale. Higher amount of B2O3 (in liquid state) in the B2O3-SiO2 scale is expected to enhance its self-healing ability. It has been further reported that the addition of La as an additive to the ZrB2-ZrC composite processed by spark plasma sintering, significantly enhances the amount of densification by forming between the powder particles a La-enriched  ⁎ Corresponding author. E-mail address: rahul@metal.iitkgp.ac.in (R. Mitra).  https://doi.org/10.1016/j.jeurceramsoc.2019.03.043 Received 28 October 2018; Received in revised form 15 March 2019; Accepted 19 March 2019 Available online 20 March 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'S.K. Kashyap and R. Mitra  liquid phase, which accelerates the mass transfer through grain boundaries [15]. Zapata et al. have reported that the addition of La2O3 improves the oxidation resistance of ZrB2-SiC and HfB2-SiC composites by forming a viscous borosilicate glass (BG) [16]. This study has reported an increase in the BG viscosity due to the homogeneous dispersion of oxide products, which in turn retards the inward diffusion of oxygen through the oxide scale. Compressive creep behavior of hot pressed pure ZrB2 and ZrB2 4 wt % Ni has been evaluated by Martinez et al. [17] in the argon atmosphere at temperatures in the range of 1400 °C-1600 °C, and the results have shown negligible creep strain in ZrB2 at 1400 °C. The reported stress exponents are 1.7 at 1500 °C, and 0.6 at 1600 °C for the tests carried out under stresses below 220 MPa. However, the Ni-doped ZrB2 has shown tragic failure at stresses above 25 MPa due to softening of the low-melting intergranular Ni-rich phase. Spivak et al. have investigated the flexural creep behavior of the hot pressed ZrB2-(20-75 mol%) ZrN composites at temperatures between 2000 °C and 2300 °C, when subjected to a stress-range of 5 MPa −20 MPa in the helium atmosphere. In this study, the highest creep rate has been observed for the composite having the composition of ZrB2-50 mol% ZrN [18]. Furthermore, Kats et al. have reported about the flexural creep behavior of two-phase ceramic composites including ZrC-ZrB2 and TiC-TiB2 systems at temperatures between 1700 °C and 2420 °C and stresses of 5-30 MPa in helium atmosphere. In the ZrC (20-79 mol%) - ZrB2 system, the creep rate has exceeded that observed for their individual constituents by one-to-two orders of magnitude [19]. A study by Mallik et al. has shown that Si3N4 addition to hot-pressed ZrB2-SiC composite lowers the compressive creep resistance for the tests carried out under 93-140 MPa at temperatures in the range of 1300 °C-1425 °C in laboratory air [20]. With the increase in test temperature, the stress exponent has decreased from 1.7 to 1.1 for ZrB2-SiC composite, and from 1.6 to 0.6 for ZrB2-SiC-Si3N4 composite. The activation energies for these composites have been found as ≈95 ± 32 kJ/mol at temperatures ≤1350 °C, and as ≈470 ± 20 kJ/ mol in the range of 1350-1425 °C. The study of post-creep microstructures and analysis of activation energies has shown that damage during creep of these composites occurs by grain boundary sliding, which is controlled by O2− diffusion through SiO2 at ≤1350 °C, and by the viscoplastic flow of the glassy interfacial film at temperatures ≥1350 °C. Observation of stress exponent values < 1 in case of the ZrB2-SiC-Si3N4 composites creep-tested at ≥1400 °C have been ascribed to grain boundary sliding involving solution-precipitation type mechanism with formation of Si2N2O at particle-matrix interfaces. A study on ZrB2-(0-50) vol% SiC composite with SiC particle sizes of 2 or 10 μm was carried out by Talmy et al. at temperatures in the range of 1200 °C-1500 °C under the applied stress of 30 to 180 MPa [21]. The results of this study have shown an increase in creep rate with increasing temperature, SiC content, stress and with decreasing SiC particle size. The stress exponent has been found as 1 for the ZrB2-(0-25) vol% SiC composite and 2 for the ZrB2-50 vol% SiC composite, which indicates the involvement of diffusional creep and grain boundary sliding as operating mechanisms, respectively. A study on creep behavior of ZrB2-SiC composites in the inert atmosphere has shown stress exponent of unity at 1400 °C, indicating the role of diffusion through ZrB2 lattice or grain boundaries or ZrB2-SiC interfaces [22]. In contrast, the stress exponents are reported to be in the range of 1.7-2.2 at higher temperatures up to 1600-1820 °C, which has indicated the role of cavitation in accommodation of grain boundary sliding. Reports on the effect of LaB6 addition on oxidation resistance of ZrB2 based composites are available in the published literature [23-26]. According to these studies, the addition of LaB6 significantly improves the oxidation resistance of ZrB2-SiC composites exposed at 2400 °C for a dwell time of 600 s and at 1600 °C for 1 h (3600 s). Such improvement has been attributed to the formation of a refractory oxide scale comprising lanthanum zirconate (La2Zr2O7) and lanthanum silicate (La2Si2O7), along with zirconia and silica [23,24]. On the other hand,  Journal of the European Ceramic Society 39 (2019) 2782-2793  some other studies have shown that the presence of LaB6 lowers the oxidation resistance of the ZrB2 based composites by lowering the eutectic temperature of oxide scale, and by increasing the concentration of oxygen vacancies [25,26]. Survey of the literature indicates that most of the previous work has focused on the oxidation behavior of the LaB6 containing composites, whereas their creep behavior remains unexplored. Considering that oxidation and creep at high temperatures are interrelated, the study of creep behavior of ZrB2-SiC-LaB6 composites in the air is of interest. In the present study, the ZrB2-SiC composite samples with varying amounts of LaB6 have been prepared by spark plasma sintering, and thereafter study of their bulk densities, microstructures, hardness, and creep behavior with due emphasis on oxide scale formation have been carried out. The creep tests have been carried out at 1300 °C (≈0.45TM, where TM = absolute melting point of ZrB2), as earlier studies have confirmed the formation of protective borosilicate scale on ZrB2-20 vol.% SiC composite subjected to isothermal exposure at this temperature in air for 24 h [27] or during creep [20].  2. Experimental procedure  The raw materials used to prepare composites for this study were ZrB2 (H.C. Starck, Germany, 5.4 ± 2.2 μm), SiC (H.C. Starck, Germany, 4.0 ± 1.4 μm), LaB6 (Alfa Aesar, Massachusetts, US, 7.9 ± 2.8 μm), and B4C (Boron Carbide India Ltd, Mumbai, India, 1.84 ± 1.3 μm) with 99.5% purity as well as phenolic resin (ABR Organics Limited, Hyderabad, India). Phenolic resin (phenol formaldehyde with ethanol as solvent), type ABRON PR100, has been used as a source of carbon in the raw materials. The carbon content of the phenolic resin has been found to be 37 wt% by thermo-gravimetric analysis. Three types of ZrB2-SiC based composites were prepared by varying their LaB6 contents as 7 vol%, 10 vol%, and 14 vol%, and these are referred to as ZSBCL-7, ZSBCL-10 and ZSBCL-14, respectively. The other constituents of the aforementioned ZrB2-based composites were SiC (20 vol%), B4C (5.6 vol%), and C (4.8 vol%). Mixing of powders was carried out using WC-Co balls and vials inside a Planetary Mono-mill (Model Pulverisette 6, Fritsch, Germany) being operated at 250 rpm for 6 h. Consolidation of the aforementioned composite powders was carried out by using spark plasma sintering (Dr. Sinter, Japan) for 5 min under a pressure of 70 MPa at 1800 °C. For this purpose, the powder sample was heated first at the rate of 70 °C/min up to 1500 °C, and then at 50 °C/min up to 1800 °C. A ram pressure of 70 MPa was applied before the onset of temperature ramp, and it was kept constant throughout the heating schedule up to the sintering temperature. During the SPS experiments, the change in the powder bed height with time was determined from the piston displacement. The value of ram displacement was subsequently corrected by subtracting the thermal expansion of the graphite die, which was calculated by heating an empty graphite die under similar experimental condition. Spark plasma sintering is reported to be a very efficient process for sintering of the UHTCs, since heating and compaction by applying load are carried out simultaneously, which not only enhances densification but also restricts grain growth due to exposure at high temperature for a short duration [28,29]. The bulk densities of the investigated composites were measured with the help of Archimedes principle using the water as medium. Different phases present in both as-sintered and post-creep samples were analyzed by X-ray diffraction (XRD) (BRUKER D8 ADVANCE) technique using Cu-Kα radiation. Microstructural examination of the investigated composites was carried out by using secondary electron (SE) and back-scattered electron (BSE) imaging modes on a field emission scanning electron microscope (FESEM, Zeiss Supra 40, Carl Zeiss NTS GmbH, Oberkochen, Germany), coupled with energy dispersive spectroscopy (EDS, Oxford Instruments, Wycombe, UK) for compositional analysis. The average grain size of the ZrB2 matrix was measured through image analysis of the SEM micrographs by the area count method involving a minimum of 200 grains. The selected polished samples were etched for 4 min by using a solution having HF,  2783  \\x0c', '3.2. Densification  Bulk densities of the composites were experimentally measured by Archimedes principle, whereas the corresponding theoretical densities were estimated using the rule of mixture (ROM) by considering the densities of ZrB2, SiC, LaB6, B4C and C, as in starting powders. The theoretical density may be considered as approximate due to the varying amounts of LaB6, B4C and C being consumed in reduction of impurity oxides during sintering of each type of composite. The relative density of each composite was calculated from the ratio of their bulk density to theoretical density. Table 1 presents the relative densities of the composites processed by SPS at 1800 °C, indicating its increase with increasing LaB6 content of the composites. Further, the plots depicting the variation of instantaneous relative densities of the composites as a function of temperature during SPS as illustrated in Fig. 2, show evidence of accelerated densification at temperatures ≥1200 °C. Here, the instantaneous relative density, D at a given temperature is related to instantaneous sample height, L as:  D = (Lf/L).Df  (1)  where Lf is the final height and Df is the final relative density [30]. The results in these plots suggest that the relative density of composites increases with increasing volume fraction of LaB6.  3.3. Phase identification by X-Ray diffraction  The XRD patterns of the composites prepared through SPS at 1800 °C as shown in Fig. 3(a), confirm the presence of all the phases except B4C and LaB6. The absence of peaks representing B4C, WC and LaB6 is attributed their low volume fraction, which indicates that much of these phases is consumed in the reduction of surface oxides. Interestingly, the XRD patterns from ZSBCL-10 and ZSBCL-14 composites show the peaks of LaBO3  HNO3, and H2O in a ratio of 2:3:95. For the study of microstructures using the transmission electron microscope (TEM), discs with 3 mm diameter were cut from thin sections of the composite samples with the help of an ultrasonic disc cutter. These samples were subsequently thinned by mechanical polishing on diamond-coated discs and dimplegrinder, followed by argon ion-thinning on a precision ion polishing system (Gatan Inc., Pleasanton, CA, USA) to electron transparency. These samples were examined using both bright-field and dark-field imaging as well as selected area electron diffraction (SAED) on a TEM (Model JEM 2100, JEOL, Tachikawa, Tokyo, Japan) operated at an acceleration voltage of 200 kV. The hardness of the as-sintered samples was measured by using a hardness tester (Model VM50, Fuel instruments & engineers Pvt. Ltd., Kolhapur, Maharashtra, India) operated using a load of 10 kgf for 15 s dwell time. Compression creep tests were performed at 1300 °C by a lever arm creep testing machine (Model 2390 Series, Applied Test System Inc., Butler, PA, USA) with 10:1 lever-arm ratio at three different loads of 75 N, 100 N and 125 N (approximately, 47 MPa, 62 MPa and 78 MPa) on all three types of investigated composites in laboratory air. The dimensions of the creep specimens were 6 mm × 4 mm × 4 mm. The specimens were polished using both coarse (6 μm), and fine (1 μm and 0.25 μm) size diamond pastes smeared on a lapping cloth to make the sample surfaces flat and smooth to avoid stress concentration. The change in specimen length during compression creep tests was measured by a linear variable differential capacitor (LVDC) having a least count of 2.54 × 10−4 mm (10-5 inch), which suggests that the minimum measurable strain is 4.2 × 10-5. Creep test was performed on each type of composite samples at a given load until the steady-state stage was reached, to measure the minimum or steady state creep rate. Subsequently, the oxide scales and cracks formed in the creep-tested samples were examined using XRD, SEM, and TEM to understand the combined effects of environment and the applied load at 1300 °C on the evolution of damage in the investigated composites.  3. Results and discussion  3.1. Raw material characterization  The XRD patterns obtained from the ball-milled powder samples, as depicted in Fig. 1(a) show the peaks of WC along with those of ZrB2, SiC, and LaB6 for all types of composites. The peaks of WC might have originated from contamination by the abrasion of WC balls and vials by SiC particles during the ball milling process. The results of Rietveld analysis have shown the presence of 1.2 wt% WC in the ball-milled starting powder mixtures used for fabrication of all types of composites investigated in this study. Due to its relatively low volume fraction, B4C peak was not visible in the XRD pattern from the as-milled powder. Examination of the SEM images [Fig. 1(b)] showing milled powder particle morphology indicates irregular shapes with sharp corners and wide size distributions. Such sharp corners in the milled powder particles are considered as responsible for abrasion of the WC vials and balls.  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Table 1 Relative densities, grain sizes and hardness of the ZrB2-SiC-LaB6 composites processed by spark plasma sintering. Here, ρth, ρSPS, and ρRELATIVE are theoretical density, density obtained by SPS, and relative density, respectively.  Composite  ZSBCL-7 ZSBCL-10 ZSBCL-14  ρth (g. cm−3)  ρSPS (g. cm−3)  ρRELATIVE (%)  ZrB2 grain size (μm)  Hardness (HV) at 10 kgf  5.03 4.99 4.93  4.78 4.83 4.85  95.0 97.5 98.2  2.6 ± 0.5 3.8 ± 1.1 1.4 ± 0.5  1400 ± 90 1500 ± 100 2040 ± 120  confirmed on examination of the bright field TEM image of ZSBCL-10 composite as shown in Fig. 7(d) and the corresponding SAED pattern in Fig. 7(e). TEM studies have also shown amorphous glassy phase to be present at some of the grain boundary triple point junctions as shown in Fig. 7(f). The amorphous structure of this phase is confirmed by diffuse halo in the SAED pattern, as shown in Fig. 7(g). The EDS analysis of this amorphous phase has shown the presence of B, C, O and Si, which indicates the presence of B2O3 and SiO2 along with unreacted carbon. A well-developed dislocation substructure has also been demonstrated by Kim and Shim indicating localized plastic deformation of the ZrB2-ZrC composite during the SPS process [15]. The dislocations in ZrB2 processed by SPS using TaSi2 as additive have been found to be straight and originating from grain boundaries, which do not exhibit any evidence of amorphous phase [33]. Therefore, the presence of high dislocation density in the ZrB2 grains of the ZSBCL-14 suggests that a significant amount of plastic deformation accompanies the process of densification by SPS due to the application of ram pressure along with heating. On the other hand, the presence of glassy phase at grainboundaries and particle-matrix interfaces of the ZSBCL-10 probably does not trigger much dislocation activity, as has been inferred in Ref. [33]. The presence of glassy phase at the matrix grain boundaries or ZrB2-SiC interfaces is expected to have enhanced intergranular diffusion causing grain growth. Therefore, it is appropriate to infer that lack of sufficient plastic deformation, as well as the presence of glassy phase at grain boundaries of the ZSBCL-10, are responsible for lower densification and larger grain size of the ZSBCL-10 composite compared to that of the ZSBCL-14.  3.5. Role of sintering additives and LaB6 on densification and microstructural evolution  Observation of XRD patterns indicates that much of B4C, LaB6 and WC is consumed in the reduction of surface oxides, i.e., ZrO2, B2O3 and SiO2 through the Reactions (1-6), as shown in Table 2. The Reactions (1-5) in this table show the effect of B4C and C during densification [34-36], whereas the Reaction (6) describes the role of LaB6 as a typical oxygen scavenger. The negative values obtained for the free energy change for each of the Reactions (1-5) as shown in Table 2, confirms their thermodynamic feasibility. Owing to the unavailability of data for LaB6 and LaBO3, it was not possible to calculate the free energy change for the Reaction (6). However, the results of an earlier study have  Fig. 2. Plots depicting the variation of instantaneous function of temperature during the SPS process.  relative density  as  a  surface topographies of the sintered composites are shown in Fig. 5(a) and (b), respectively. It is evident from these figures that with increase in the LaB6 content, the volume fraction of porosity is decreased, which is consistent with the results of density measurements. Although LaB6 is not detectable in the microstructures of the sintered composites (Fig. 4), yet the combination of SEM (BSE) image shown in Fig. 6(a) along with accompanying EDS X-ray maps of La, B, and O shown in Fig. 6(b), confirms the formation of LaBO3 at a triple junction between the grains of ZrB2, SiC, and B4C. The average sizes of the ZrB2 matrix grains in the investigated composites obtained through measurement of sizes of ≥200 grains by analyzing several images depicting the microstructures are shown in Table 1. It is interesting to note that the average size of the ZrB2 grains in the ZSBCL-14 composite is finer than that in other two composites, which suggests that grain coarsening during sintering of the composites is partially restricted by LaB6 addition. The bright field TEM images of the ZSBCL-14 composite depict the grains of ZrB2 and SiC in Fig. 7(a) and (b), as confirmed by both selected area diffraction and EDS analysis. The bright field TEM images as shown in Fig. 7(a) and (b) indicate the presence of high dislocation density in the ZrB2 grains of the sintered composite. The identity of ZrB2 grains is confirmed by the SAED analysis, as shown in Fig. 7(c). Furthermore, the presence of LaBO3 at multigrain junctions is  Fig. 3. a) XRD patterns showing the peaks of composites processed by spark plasma sintering at 1800 °C; and (b) enlarged view of LaBO3 peaks.  2785  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 4. SEM (BSE) images depicting the microstructures of (a) ZSBCL-7, (b) ZSBCL-10, and (c) ZSBCL-14 composites.  Fig. 5. SEM (SE) images showing the surface topographies of (a) ZSBCL-7 and (b) ZSBCL-14 composites. The pores are shown with arrows.  Fig. 6. The SEM (BSE) image of the ZSBCL-10 composite, showing (a) the presence of LaBO3 and (b) a typical inter-grain junction between ZrB2, SiC and B4C; as well as elemental EDS maps confirming the presence of LaBO3.  shown that Reaction (6) is expected well below the sintering temperature [31]. Indeed, the observation of LaBO3 next to B4C at an intergrain junction in the microstructure (Fig. 6(b)) confirms the occurrence of Reaction (6). However, LaBO3 is not always found adjacent to B4C as shown in Fig. 6(a) and 7(d), which suggests its formation by oxidation of LaB6 during sintering due to the presence of oxygen, which is probably trapped as impurity during ball-milling and subsequent powder handling. The reactions mentioned in Table 2 show how the addition of B4C, C, and LaB6 aids in densification of the investigated ZrB2-based composite during SPS by reducing the oxides at their surfaces. The increase in relative density with LaB6 content as shown in Table 1 suggests that its oxygen scavenging plays a key role in the process of densification. Further, the ZSBCL-14 being processed with the highest LaB6 addition also exhibits the finest matrix grain size, as the absence of glassy intergranular film restricted grain boundary diffusion and accompanying grain growth during the short duration of SPS processing. On the other hand, the presence of intergranular glassy phase in the ZSBCL-10 as shown in Fig. 7(d-g) has led to enhanced grain boundary diffusion and grain growth during densification by SPS.  3.6. Hardness  From the results shown in Table 1, it is clear that the hardness value of the ZSBCL-10 composite is only marginally greater than that of the ZSBCL-7. However, the hardness of the ZSBCL-14 is higher than that of the ZSBCL-10 composite by 36%. It may be noted that the ZSBCL-10 has not only much greater relative density, but also coarser grain size compared to that of the ZSBCL-7. Whereas hardness of the composites is expected to scale with their relative density, it is lowered with increase in matrix grain size. Therefore, the effect of higher relative density of ZSBCL-10 appears to be neutralized by its coarser matrix grain size. Furthermore, the ZSBCL-14 composite with both higher relative density and finer grain size, exhibits higher hardness than that of ZSBCL-7 or ZSBCL-10.  3.7. Compression creep behaviour  3.7.1. Creep results The variation of  creep strain rate with applied stress at a given  2786  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 7. Bright field TEM images of the ZSBCL-14 composite showing the (a) interface between ZrB2 and SiC grains, and (b) dislocation tangles inside the ZrB2 grain; with (c) typical SAED pattern from a ZrB2 grain. Microstructure of the ZSBCL-10 composite: (d) bright field TEM image showing the presence of LaBO3 at inter-grain junction, and (e) corresponding SAED pattern; (f) bright field TEM image showing the presence of amorphous phase at inter-grain junction, and (g) corresponding SAED pattern showing diffuse halo confirming the presence of amorphous phase.  temperature is expressed using the following relation:  ε´=Aσn  (2)  where ε´ is the strain rate, A is a constant, and n is stress exponent. Creep tests were performed at 1300 °C (≈0.45TM, where TM = absolute melting point of ZrB2) under three different stresses, that is, 47 MPa, 62 MPa, and 78 MPa. The results obtained for compression creep tests carried out at 1300 °C under different stresses on ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites are respectively, shown in Fig. 8(a), (b), and (c). Plots depicting the variation of strain with time for the ZSBCL-7 composite creeptested at 47 MPa shows that steady state is reached after 5 h, and the corresponding strain rate (ε´) as found from the slope of the best-fit line is found as 7.0 × 10−8 s-1. It may be noted that at the higher stresses of 62 MPa and 78 MPa, the ZSBCL-7 samples were disintegrated prematurely, and therefore further tests at higher loads were not carried out on this composite. The log-log plots of strain-rate against stress as shown in Fig. 8(d), indicate that the steady-state strain rates of the ZSBCL-14 under different stresses are higher than those of the ZSBCL-10 by nearly 3 times. Furthermore, the values of n determined from the slopes of best-fit lines for log-log plots of steady-state creep rate against stress in this figure are found to be 1.3 ± 0.1 and 2.6 ± 0.2 for ZSBCL-10 and ZSBCL-14, respectively. These values of stress exponents are suggestive of diffusioncontrolled (expected n1) and dislocation-glide controlled (expected n3) for mechanisms to be operative for creep deformation of ZSBCL-10 and ZSBCL-14 composites, respectively.  3.7.2. Study of post-creep microstructures and oxide scale formed on creeptested specimens X-ray diffraction analysis of the oxide scales formed on the samples creep-tested at 1300 °C has shown the presence of monoclinic La2(Si2O7), tetragonal ZrSiO4, tetragonal ZrO2, and monoclinic ZrO2, as shown in Fig. 9. In the XRD pattern from the oxide scale of the ZSBCL-7 composite as shown in Fig. 9(a), very few peaks of La2Si2O7 are observed due to low volume fraction of LaB6 in this composite. In contrast, the higher intensity peaks of La2Si2O7 in Fig. 9(c) indicate a higher amount of its formation in the protective oxide scale due to greater volume fraction of LaB6 in the ZCBCL-14 composite. By comparing Fig. 9(a) through (c), it is also evident that the peaks of tetragonal ZrO2 also exhibit increase in height with increasing amount of LaB6. This observation is consistent with an earlier report by Eakins et. al showing that the presence of LaB6 restricts the tetragonal to monoclinic transformation of ZrO2 [37]. The tetragonal to monoclinic transformation of ZrO2 is known to cause volume expansion, and generate compressive stress, which in turn leads to buckling of thin oxide scale followed by its cracking. Therefore, the inhibition to tetragonal to monoclinic transformation of ZrO2 is expected to have contributed to the structural stability of the protective oxide scale. Furthermore, microstructural investigations of the post-creep composites using SEM accompanied by EDS analysis have also confirmed the presence of the aforementioned oxides on the surfaces of the composites, as is evident from the results depicted in Fig. 10.  Table 2 Reactions involving B4C, C, and LaB6 for reduction of surface oxides along with free energy change at 2100 K (1827 °C).  Reaction No.  Reaction  1. 2. 3. 4. 5. 6.  5B4C (s) + 7ZrO2 (s) → 7ZrB2 (s) + 5CO (g) + 3B2O3 (l) B2O3(l) + ZrO2(s) + 5C (s) → ZrB2 (s) + 5CO (g) 2B2O3 (l) + 7C(s) → B4C (s) + 6CO (g) SiO2 (s) + 3C → SiC (s) + 2CO (g) 7SiC (s) + 6B2O3 (l) → 7SiO2 (s) + 3B4C (s) + 4CO (g) 4LaB6 (s) + 4B2O3 (l) + 7C (s) → 4LaBO3 (s) + 7B4C (s)  2787  Free Energy Change (kJ/mol)  −514 −252 −250 −101 −41 Data not available  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 8. Typical plots of creep strain against time for (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites; as well as (d) log-log plots of strain rate against stress for ZSBCL10 and ZSBCL-14 composites.  Examination of the microstructures of the post-creep ZSBCL-14 composite samples has shown the presence of micro-cracks at ZrB2 matrix grain boundaries and ZrB2-SiC interfaces, as shown in Fig. 11(a), thereby suggesting that grain boundary sliding is the primary mechanism for creep damage. Chemical analysis of the cross section of the oxide scale of the above composite by EDS analysis, as shown in Fig. 11(b), depicts the presence of lanthanum rich phase formed at the outermost surface, which may be the combination of both La2Si2O7 and La2Zr2O7 phases, followed by layers typically comprising ZrSiO4, a mixture of ZrO2+SiO2+SiC, and SiO2. Since lanthanum is more reactive than other elements in the composite, lanthanum hexaboride is preferentially oxidized, and forms a layer comprising La2(Si2O7) and La2Zr2O7 on the sample at an early stage of oxidation. The oxygen  anions are expected to diffuse through this outer layer to react with the parent composite to form other phases at the oxide-ceramic interfaces until the outer layer comprising La2(Si2O7) and La2Zr2O7 becomes continuous. In tune with the observations recorded for the ZSBCL-14 in Fig. 11(a), Fig. 12(a) and (b) depicting the microstructures of the bulk portion of post-creep ZSBCL-10 composite at same scale also show the presence of interfacial microcracks at ZrB2-SiC boundaries. Furthermore, investigation of the cross-sectional microstructure depicting the oxide scale formed on the post-creep ZSBCL-10 composite reveals the presence of interfacial cracks, as shown in Fig. 12(c). However, the isothermally exposed samples show the presence of crack-free and more adherent oxide scale as shown in Fig. 12(d). This figure shows that the  Fig. 9. XRD patterns from the oxide scales of: (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites.  2788  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 10. SEM (SE) images depicting the oxide scales of (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites, subjected to creep tests.  oxide scale contains ZrSiO4 as the outer layer, with ZrO2 and B2O3-SiO2 forming beneath. From the SEM (BSE) image of the oxide scale formed by isothermal exposure as shown in Fig. 12(b), it is apparent that B2O3SiO2 aids in sintering or cementing leading respectively, to the closure of pores or cracks in the oxide scale. However, the B2O3-SiO2 containing locations in the oxide scale of the creep-tested sample (Fig. 12(a)) appear to be cracked, indicating the absence of its cementing action under an externally applied load. The chemical analysis of a typical oxide scale cross-section on the post-creep ZSBCL-10 composite has been further investigated by EDS elemental mapping, and the results are shown as SEM (BSE) image along with EDS X-ray maps in Fig. 13. Examination of the EDS X-ray maps shows the outermost part of the oxide scale as primarily consisting of La and O along with some amount of Zr and Si. Followed by the aforementioned La-rich outer layer, enrichment of Zr, O, and Si occurs, which confirms the presence of ZrSiO4. The results obtained from the EDS mapping are consistent with those obtained by XRD (Fig. 9) showing the presence of La2Zr2O7 and La2Si2O7 at the surface followed by ZrSiO4. The EDS map of Si in Fig. 13(c) shows the presence of a SiC-depleted region at the compositeoxide interface, adjacent to which is a more or less continuous SiO2-rich layer as one proceeds towards the outer surface, which suggests that SiC undergoes selective oxidation. By comparing the results in Figs. 12 and 13 with that in Fig. 11(b), it is evident that the oxide scale on the ZSBCL-10 does not have a continuous outer layer containing a mixture of La2Zr2O7 and La2Si2O7, unlike that in case of the ZSBC-14 composite. Table 3 shows the values of the thickness of oxide scales formed on the post-creep samples and those exposed isothermally at 1300 °C for similar durations without any load application. The results shown in this table indicate that the oxide scale thickness increases with test duration as expected, but remains < 10% of the specimen width even  after exposure for the entire duration of creep test. Further, it is also possible to see that that the average thickness of oxide scale formed on a creep-tested sample is higher than that grown on isothermal exposure for similar conditions of time and temperature without load. This observation may be because of the damage caused in the oxide scale during creep under compressive stress, which probably accelerates oxidation of the specimen. Quantitative image analysis has shown the net length of cracks per unit area to be similar, i.e., in the range of 3-4% in the creep-tested samples of both ZSBCL-10 and ZSBCL-14 composites. However, a comparison of the oxide-scale thickness values as shown in Table 3 indicates that the sum of the average growth rate per hour of the oxide scale and depth of oxygen penetration during creep is 26.7 μm/h for ZSBCL-7, 27.6 μm/h for ZSBCL-10 and 20.8 μm/h for ZSBCL-14. It is also noted that the sum of the average growth rate per hour of oxide scale and depth of oxygen penetration during creep in the ZSBCL-10 composite is greater than that in case of the ZSBCL-14 by 20.5% and 46.9%, respectively during exposure to creep. In contrast, the average growth rate of the oxide scale of the ZSBCL-10 composite during isothermal exposure at 1300 °C is comparable to that of the oxide scale formed on the ZSBCL-14. These observations suggest that the oxidation resistance of the ZSBCL-14 composite is superior to that of the ZSBCL-10 during exposure to creep. On the other hand, the oxidation resistances of these two composites are almost comparable for isothermal exposure at this temperature. Lower oxygen penetration in the ZSBCL-14 compared to that in the ZSBCL-10 during creep indicates that the oxide scale is more adherent and protective in case of the former composite, probably due to the formation of La2Zr2O7 + La2Si2O7 as the outer layer. Previous studies on ZrB2-SiC-LaB6 composites have shown that the formation of a refractory oxide product like lanthanum zirconate  Fig. 11. SEM (BSE) images depicting the microstructures of the ZSBCL-14 composite samples creep-tested at 1300 °C under 78 MPa, showing (a) microcracks at grain boundaries, and (b) constituents of various layers in the oxide scale.  2789  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 12. SEM (BSE) images of the ZSBCL-10 composite: Post-creep microstructure in the sample tested at 1300 °C under 78 MPa, showing (a) creep cavitation, (b) interfacial cracks at ZrB2-SiC boundaries; as well as the cross-sectional microstructures of the oxide scales formed in the sample (c) creep-tested at 1300 °C under 47 MPa for 8 h, and (d) exposed isothermally at 1300 °C for 8 h without any load.  (La2Zr2O7), which is reported as stable up to 2400 °C, is capable of retarding the inward diffusion of oxygen anions [23]. Therefore, it is appropriate to infer that formation of lanthanum zirconate in the oxide scale of the ZSBCL-14 composite during creep exposure has hindered further oxidation of the composite. In addition, lanthanum silicate is known to be in the form of a continuous glassy layer [26,38], which undergoes viscoplastic flow and fills pores and cracks within the oxide scale to retard the ingress of molecular oxygen. In this manner, formation of lanthanum silicate contributes to self-healing ability of the oxide scale. Bright-field TEM images of a post-creep sample of the ZSBCL-14 composite depicting the presence of dislocations in the ZrB2 grains and microcracks at the grain boundaries are shown in Fig. 14(a) and (b), respectively. The presence of dislocation arrays inside the ZrB2 grains as shown in Fig. 14(a) is suggestive of their plastic deformation due to the pressure applied during either SPS or creep tests. Furthermore, a bright field TEM image of the creep-tested ZSBCL-10 composite showing the amorphous phase at grain boundary junction along with representative SAED pattern is shown in Fig. 14(c). This amorphous phase appears to be similar to that observed in the pre-creep microstructures of the ZSBCL-10 composite, and may not have formed by percolation of oxygen through oxide scale during creep. Furthermore, the bright-field TEM images as shown in Fig. 14(d) and (e) respectively, depict the presence of dislocations in ZrB2 grains and microcracks along the ZrB2 grain boundaries or ZrB2-SiC interfaces. These images suggest that plastic deformation and grain boundary or interfacial cavitation are the mechanisms operating during creep. In comparison to the post-creep microstructures of the ZSBCL-14, the dislocation density is found to be relatively less in the ZrB2 grains of the ZSBCL-10, which again may be attributed to the absence of dislocations or low dislocation density in its pre-creep microstructure. The presence of grain boundary microcracks in the TEM images as shown in Fig. 14(b) and (e), confirms the role of grain boundary sliding as the mechanism of creep damage.  3.7.3. Mechanism of creep: Role of LaB6, dislocations and grain size The results of microstructural studies have shown that the presence of LaB6 causes scavenging of oxygen impurities, thereby aiding in densification and inhibition of grain growth. Incomplete densification and damage by accelerated oxidation are probably the reasons for the disintegration of ZSBCL-7 at higher stresses. It is intuitive that oxidation would affect the creep behaviour in air at high temperature, within the depth of the sample being penetrated by oxygen. Penetration of oxygen and formation of glassy films at ZrB2 grain boundaries and ZrB2-SiC  interfaces would worsen the creep resistance by promoting grain boundary sliding, which in turn leads to damage by intergranular or interfacial cracking as shown in Fig. 12(c). Therefore, the formation of an outer layer of La2Zr2O7 and La2Si2O7 as well as an inner layer of glassy B2O3-SiO2 in the oxide scale of the creep-tested samples in the high LaB6 containing composite samples (ZSBCL-14) may be considered as desirable to restrict internal oxidation of grain boundaries. It is interesting to note that the stress exponent of ZSBCL-14 (n = 2.6 ± 0.21) is two times higher than that of ZSBCL-10 (n = 1.3 ± 0.08), in spite of the finer grain size of the former composite. A stress exponent close to 3 in case of the ZSBCL-14 composite is suggestive of glide controlled dislocation creep being operative, as is evident from the TEM image in Fig. 14(a). Furthermore, TEM examination of the pre-creep tested ZSBCL-14 composite has shown a high density of dislocations in many of the ZrB2 grains, and their formation is attributed to pressure applied during SPS at 1800 °C [Fig. 7(b)]. It is intuitive that glide and multiplication of these dislocations with the transfer of slip to neighboring less favorably oriented (for deformation) matrix grains, is operative during creep of the ZSBCL14 composite. Earlier studies have shown the prevalence of prismatic and pyramidal slip in the composites subjected to a flexural test at 1000 °C, and that of basal slip in the samples tested at 1500 °C [39]. The metallic nature of chemical bond is said to be responsible for dislocation activity in ZrB2 even on being subjected to indentation or scratch at the ambient temperature [40]. The presence of long and straight dislocations in ZrB2 grains of the creep-tested samples as shown in Fig. 14(a), may be attributed to the Peierls barrier expected in ceramic crystals with covalent or ionic bonding. Furthermore, the absence of five independent slip systems in hexagonal structured ZrB2 is expected to create intergranular or interfacial discontinuities or cracks unless the resulting strain mismatch is accommodated by grain boundary diffusion, as is expected for a fine-grained matrix of the ZSBCL-14 composite. Therefore, it is proposed that the ease of deformation of the ZSBCL14 by dislocation mechanism is responsible for its higher steady-state creep rates compared to those observed for the ZSBCL-10, which shows much lower density of dislocations in its post-creep microstructure. The absence of dislocations in pre-creep microstructures of the ZSBCL-10 is probably responsible for a mechanism different from that of the ZSBCL14. This observation may be related to the presence of glassy phase at the inter-grain boundaries in the pre-creep ZSBCL-10 composite, considering that the earlier study by Hu et al. has reported dislocation activity only in the ZrB2-SiC composite having clean grain boundaries, with no glassy phase [33].  2790  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 13. Microstructure of the ZSBCL-10 composite creep-tested at 1300 °C under 78 MPa showing (a) SEM (BSE) corresponding EDS X-ray maps showing the areas enriched in (b) Zr, (c) Si, (d) La, (e) O, (f) B, and (g) C.  image depicting the oxide scale; as well as  Table 3 The oxide scale thickness of the samples creep-tested under 47 MPa stress and isothermally exposed for same duration without stress at 1300 °C.  Composites  Creep/ Oxidation duration (h)  Oxide scale thickness toxide (μm) and O2 penetration  Percentage of oxide scale thickness 2.toxide/tspecimen (%)  ZSBCL-7 ZSBCL-10 ZSBCL-14  6 8 12  Creep  Oxide  Oxidation  Oxygen penetration  Oxide  Oxygen penetration  67 ± 11 141 ± 21 175 ± 18  93 ± 15 80 ± 14 75 ± 8  60 ± 16 89 ± 31 145 ± 25  77 ± 11 51 ± 23 34 ± 9  Creep  3.4 7.1 8.7  Oxidation  3.0 4.5 7.2  Grain boundary sliding leading to the formation of microcracks at grain boundaries or ZrB2-SiC interfaces appears to be the major mechanism of damage in both ZSBCL-10 and ZSBCL-14, irrespective of the dominance of dislocation creep or diffusional creep. It is a well-known fact that intergranular strain introduced by grain boundary sliding during creep needs to be accommodated by either diffusion or slip involving the formation of geometrically necessary dislocations to avoid  intergranular or interfacial cracking. In an earlier study on the creep of the ZrB2-SiC composite at 1800 °C, grain boundary sliding accommodated by dislocation activity has been reported as the mechanism for large creep strains [41]. Formation of geometrically necessary dislocations in ZrB2 matrix due to strain gradients near grain boundaries and their role in accommodation of grain-boundary sliding strains as the rate-controlling mechanism has been noticed in this study. Role of  2791  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  Fig. 14. Bright field TEM images of the postcreep ZSBCL-14 composite depicting (a) dislocation array and dislocation pile-up in ZrB2 grains, as well as (b) microcrack formation at the grain boundaries. Bright field TEM images depicting the post-creep microstructures of the ZSBCL-10 composite showing (c) the amorphous phase being confirmed from diffuse halo in the SAED pattern shown as inset, (d) dislocations (arrowed) in a ZrB2 grain, and (e) microcracks (shown with arrows) formed at ZrB2 grain boundaries and ZrB2-SiC interfaces during creep.  scavenging of oxygen by B4C, C and LaB6. For the ZrB2-SiC composites having relative density ≥97%, the ZrB2 grain size decreases with increasing LaB6 content. (3) Steady-state creep rate increases with applied stress, but it remains within the range of 10−8 s-1 for all the composites. The values of stress exponents as well as examination of post-creep microstructures suggest dislocation glide controlled deformation to be effective in ZSBCL-14 (n = 2.6 ± 0.21), and diffusion controlled creep in ZSBCL-10 (n = 1.3 ± 0.08). (4) Creep resistance of the ZSBCL-10 composite is found to be superior to that of ZSBCL-14, and this is ascribed to finer grain size and easier dislocation slip in the latter composite. Whereas the presence of glassy phase at grain boundaries in the sintered ZSBCL-10 composite may be responsible for lack of dislocation slip in its ZrB2 grains during creep, its coarser grain size along with strong dependence of diffusion creep has contributed to its lower steady state strain rate. (5) Microstructural studies of the creep-tested ZSBCL-10 and ZSBCL-14 have indicated the formation of cracks at ZrB2 grain boundaries as well as at ZrB2-SiC interfaces, which suggest that grain-boundary sliding is the major operating mechanism. The strain incompatibility caused by sliding between neighbouring grains appears to be partially accommodated by grain boundary diffusion in the ZSBCL10 and dislocation slip in the ZSBCL-14.  Acknowledgements  Technical assistance received from Mr. Srikrishna Maity, Mr. Mithun Das, Mr. B. Santu Mudliyar, Mr. Suman Sarkar, and Mr. Subhabrata Chakraborty, Staff members of Central Research Facility, IIT Kharagpur for characterization of specimens, is gratefully acknowledged.  [1]  [2]  P. Kolodziej, Aerothermal Performance Constraints for Hypervelocity Small Radius Unswept Leading Edges and Nosetips, NASA Tech. Memo, 1997, p. 112204. T.A. Jackson, D.R. Eklund, A.J. Fink, High speed propulsion: performance advantage of advanced materials, J. Mater. Sci. 39 (2004) 5905-5913, https://doi. org/10.1023/B:JMSC.0000041687.37448.06. [3] D.M. Van Wie, D.G. Drewry, D.E. King, C.M. Hudson, The hypersonic environment: required operating conditions and design challenges, J. Mater. Sci. 39 (2004) 5915-5924, https://doi.org/10.1023/B:JMSC.0000041688.68135.8b.  dislocation glide and climb in accommodation of grain rotation and translation contributing to grain boundary sliding during flexural creep at 1800 °C has also been reported by Bird et al. [42]. The strain incompatibility due to unequal dislocation slip in the neighboring ZrB2 grains of the ZSBCL-14 composite has to be accommodated by grain boundary sliding. Therefore, higher steady state creep rate of the ZSBCL-14 compared to that of the ZSBCL-10 may be ascribed to not only to the ease of deformation of the ZrB2 grains in the former composite, but also the requirement of plastic strain accommodation by grain boundary sliding due to lack of five independent slip systems. As the average grain size of the ZSBCL-14 composite is nearly 1/3rd of that in the ZSBCL-10 (Table 1), the damage by grain boundary sliding is expected to be worse in the former composite. Although activation energy for creep has not been determined in the present study, yet grain boundary diffusion may be considered as the rate-controlling mechanism for the ZSBCL-10 composite considering the value of n  1 as well as the results of an earlier study on ZrB2-SiC composite with slightly coarser grain size [20]. This inference may be considered as intuitive, considering that lattice diffusivity of the covalently bonded ZrB2 and SiC is poor, and therefore its role will be relatively insignificant at 1300 °C, which corresponds to a low homologous temperature (≈0.45TM). As grain boundary diffusion is considered as the mechanism controlling the rate of creep deformation of the ZSBCL10 composite, its coarser matrix grain size compared to that of the ZSBCL-14 may be considered as responsible for its lower steady state creep rate (Fig. 8). Although the ingress of oxygen and thickness of oxide scale formed under creep condition may be marginally greater in the ZSBCL-10 than that in the ZSBCL-14 composite, yet the presence of coarser matrix grain size in the former composite appears to have lowered the steady state creep rate of the former composite.  4. Conclusions  Based on the results obtained in the present lowing conclusions can be drawn:  investigation,  the fol References  (1) The relative density and hardness values of the ZrB2-SiC based UHTC composites increase with increasing amount of LaB6. (2) Grain growth is found to be limited after SPS, which may be attributed to rapid heating and brief duration of exposure at high temperature, inhibition of grain growth by SiC, as well as  2792  \\x0c', 'S.K. Kashyap and R. Mitra  Journal of the European Ceramic Society 39 (2019) 2782-2793  [6]  [9]  [10]  [12]  [4] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. 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Soc. 90 (2007) 3660-3663, https://doi.org/10.1111/j.1551-2916.2007.01936.x. [13] H. Ma, H. Liu, J. Zhao, F. Xu, G. Zhang, Pressureless sintering, mechanical properties and oxidation behavior of ZrB2 ceramics doped with B4C, J. Eur. Ceram. Soc. 35 (2015) 2699-2705, https://doi.org/10.1016/j.jeurceramsoc.2015.03.030. [14] X. Lin, S. Ai, Y. Feng, D. Gao, X. Guo, Y. Liu, B. Xie, Fabrication and properties of insitu pressureless-sintered ZrB2 / B4C composites, Ceram. Int. 43 (2017) 15593-15596, https://doi.org/10.1016/j.ceramint.2017.08.115. [15] K.H. Kim, K.B. Shim, The effect of lanthanum on the fabrication of ZrB2-ZrC composites by spark plasma sintering, Mater. Charact. 50 (2003) 31-37, https://doi. org/10.1016/S1044-5803(03)00055-X. E. Zapata-Solvas, D.D. Jayaseelan, P.M. Brown, W.E. Lee, Effect of La2O3 addition on long-term oxidation kinetics of ZrB2-SiC and HfB2-SiC ultra-high temperature ceramics, J. Eur. Ceram. Soc. 34 (2014) 3535-3548, https://doi.org/10.1016/j. jeurceramsoc.2014.06.004. J.J. Meléndez-Martínez, A. Domínguez-Rodríguez, F. Monteverde, C. Melandri, G. De Portu, Characterisation and high temperature mechanical properties of zirconium boride-based materials, J. Eur. Ceram. Soc. 22 (2002) 2543-2549, https:// doi.org/10.1016/S0955-2219(02)00114-0. I.I. Spivak, R.A. Andrievskii, V.V. Klimenko, V.D. Lazarenko, Creep in the binary systems TiB2 -TiC and ZrB2 ZrN, Sov, Powder Metall. Met. Ceram. 8 (1974) 617-620. S.M. Kats, S.S. Ordan’yan, Compressive creep of alloy of the ZrC-ZrB2 and TiC-TiB2 systems, Sov. Powder Metall. Met. Ceram. 12 (1981) 70-75. [20] M. Mallik, K.K. Ray, R. Mitra, Effect of Si3N4 addition on compressive creep behavior of hot-pressed ZrB2-SiC composites, J. Am. Ceram. Soc. 97 (2014) 2957-2964, https://doi.org/10.1111/jace.13022. I.G. Talmy, J.A. Zaykoski, C.A. Martin, Flexural creep deformation of ZrB2/SiC ceramics in oxidizing atmosphere, J. Am. Ceram. 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  "_id": 46,
  "PDF": "Effect of moisture on the oxidation behavior of ZrB2.pdf",
  "Text": "['Received: 30 March 2020   DOI: 10.1111/jace.17500    |   Revised: 21 August 2020   |   Accepted: 14 September 2020  O R I G I N A L A R T I C L E  Effect of moisture on the oxidation behavior of ZrB2  Ravisankar\\xa0Naraparaju1 Gregory E.\\xa0Hilmas2  1German Aerospace Center (DLR), Institute  of Materials Research, Cologne, Germany 2Department of Materials Science and  Engineering, Missouri University of  Science & Technology, Rolla, MO, USA  Correspondence  Ravisankar Naraparaju, German Aerospace  Center (DLR), Institute of Materials  Research, 51170, Cologne, Germany. Email: ravisankar.naraparaju@dlr.de  |   Keyur\\xa0Maniya1   |   Alec\\xa0Murchie2   |   William G.\\xa0Fahrenholtz2   |     Abstract  Oxidation studies of ZrB2 were performed under wet air and dry air conditions at  1200°C, 1400°C, and 1500°C for 1, 4, and 10\\xa0h. Compared to dry air, the presence  of water vapor was found to enhance the oxidation kinetics by a factor of 7 to 30, depending on the temperature. Thermodynamic calculations suggested that water vapor  promoted the formation of additional volatile species such as boric acid (HBO2),  in addition to boria (B2O3) produced in dry air, which increased the evaporation  rate of B2O3. Compared to dry air, the presence of water vapor leads to more rapid  evaporation of boria and the transition from parabolic oxidation kinetic behavior (ie,  rate controlled by diffusion through boria) to linear (ie, underlying ZrB2 is directly  exposed to the oxidizing environment) at shorter times and lower temperatures.  K E Y W O R D S oxidation, ultra-high temperature ceramics, zirconia  1   |   I N T RO D U C T I O N  Ultra-high-temperature ceramics (UHTCs) are a class of  mater ials that include dibor ides (ie, ZrB2, HfB2), carbides  (ie, ZrC, HfC), and nitr ides (ie, ZrN, HfN) of early transition metals.1 UHTCs have melting  temperatures over  3000°C and the capability to tolerate extreme heating environments.2-4 Increasing interest in hypersonic vehicles  has led the development of new UHTC mater ials for wing  leading edges and nose tips, as well as propulsion system  components. Dibor ide-based UHTCs have per formance  advantages in some hypersonic applications compared to  the carbides and nitr ides due to better oxidation resistance  and ability to transfer and redistr ibute heat (thermal conductivity\\xa0>\\xa0100\\xa0W/m K at 25°C) at elevated temperatures.5  Extensive research has been conducted over the years to  understand the oxidation behavior of zirconium dibor ide  (ZrB2) in the low to high-temperature regimes. Previous   studies have divided oxidation behavior  into  three  regimes6: (a)  the  low-temperature regime below 1000°C  where a crystalline zirconia (ZrO2) and a continuous liquid/glassy bor ia (B2O3) scale form on the sur face of the  un-oxidized ZrB2 matr ix, providing passive oxidation protection. The oxidation kinetics in this regime is generally  parabolic in nature. The oxidation rate is controlled by the  diffusion of oxygen in the bor ia glass; (b) a second regime  between 1000°C and 1800°C, where the evaporation of  B2O3 begins in addition to the ZrO2 formation and the f inal  regime; and (c) a regime above 1800°C where the evaporation of B2O3 is rapid such that nearly all of the B2O3 is lost  by evaporation and a porous ZrO2 scale forms and offers  no resistance to oxygen transpor t. Bor ia volatilizes over  a wide range of conditions, such as var ious temperatures  and par tial pressures of oxygen in the external atmosphere,  and B2O3 is the most predominant volatile species in air at  1500°C.7 At temperatures above 1100°C in the presence of   This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original  work is properly cited. © 2020 The Authors. Journal of the American Ceramic Society published by Wiley Periodicals LLC on behalf of American Ceramic Society (ACERS)  J Am Ceram Soc. 2020;00:1-9.   wileyonlinelibrary.com/journal/jace  |   1          \\x0c', '2   |   water vapor, HBO2(g) was the predominant gaseous species formed by oxidation of BN, although other species  such as H3BO3 (g) and H3B3O6 (g) become impor tant at  higher water vapor contents.8 Histor ical studies on the oxidation behavior of pure ZrB2 were mainly concentrated on  understanding the effect of f low rate, oxygen par tial pressure and external pressure.9,10 In one of the studies, the rate  of oxidation of ZrB2 was found to be directly propor tional  to the oxygen par tial pressure (100-740\\xa0 Tor r) within the  temperature range of 945°C-1256°C.11 In another study  no oxygen pressure dependency at 1287°C and 1557°C  was found.10 Brown has observed an accelerated oxidation  rate in pure oxygen compared to dry air and a much higher  rate in moist air than in dry air in the temperature range  649°C-1315°C.12 The effect of water vapor on the oxidation behavior of ZrB2-SiC revealed that vapor enhances the  oxidation rate by a factor of 2-3, which was attr ibuted to  the enhanced volatilization of silica.13,14 The purpose of this study was to analyze the effect of water  vapor on the oxidation behavior of ZrB2. Thermodynamic  calculations and experimental studies of scale thickness as a  function of temperature and time were combined to identify  the underlying mechanisms.  2   |   E X P E R I M E N TA L D E TA I L S  ZrB2 specimens with relative densities of\\xa0>\\xa095% and average grain sizes of 19.3\\xa0±\\xa013.0\\xa0µm were used in the oxidation  tests. Specimens were prepared by hot pressing commercial  powders (Grade B, HC Starck) at 2100°C using 0.5\\xa0 wt%  carbon as a sintering aid. Details of the processing methods and microstructure analysis are described elsewhere.15  Three types of oxidation experiments were performed in  this study.  1. Oxidation under static ambient air (AA) 2. Oxidation under flowing synthetic air (SA) 3. Oxidation under flowing synthetic air\\xa0 +\\xa0 water vapor  (SA\\xa0+\\xa0Bubbler)  Oxidation exper iments were car r ied out in a box furnace (AA conditions, Netzsch) and a tube furnace consisting of a 1.5\\xa0m long alumina tube with an outer diameter of  10\\xa0 cm (SA conditions, Naber therm). The specimens were  heated in alumina crucibles f illed with zirconia powder to  reduce the propensity for specimens to stick to the specimen suppor t. Rectangular coupons were heated at 10°C/ min to selected temperatures, which were 1200°C, 1400°C,  and 1500°C, followed by isothermal holds of 1, 4, or 10\\xa0 h  under ambient air pressure (ie, nominally 1\\xa0atm). The coupons were cooled to room temperature at 10°C/min in the  furnace. A f lowmeter was used to control the f low rate of  synthetic air (0.6\\xa0cm/s, linear f low rate was calculated from  the volumetr ic f low rate and the diameter of the tube). For  some SA exper iments, a bubbler was used to saturate the  SA with water pr ior to f lowing into the furnace tube. The  oxidized coupons were cut, mounted in a conductive epoxy,  and polished using standard metallographic techniques to  a 0.05\\xa0 µm sur face f inish. Scanning electron microscopy  (SEM) was per formed using a DSM Ultra 55 (Carl Zeiss  NTS, Wetzlar, Germany) microscope. SEM was equipped  with an energy dispersive X-Ray spectroscopy (EDS) system (Inca, Oxford Instruments Abingdon, UK), which was  used to determine compositions of oxide phases. The oxide  scale thicknesses were measured and compared among all  the specimens. Thicknesses were always measured at the  center area of the coupon (ie, away from corners) and at  least four to f ive measurements were made for each oxidation condition. X-ray diffraction (XRD) was used to identify the crystalline phases in the oxide scales. The XRD  measurements were made using a Siemens D5000 diffractometer equipped with a CuKα radiation source with a secondary graphite monochromator (EVA/Topas 4.2 software  package, Bruker AXS, Karlsruhe, Germany) directly on  the oxidized specimens. Glow discharge optical emission  spectroscopy (GDOES) is a technique based on identifying  emissions from atoms by means from plasma by sputter ing  and GDA650 (Spectruma, Kleve, Germany) was used to  identify boron oxide in the scale. Thermodynamic calculations were per formed using the Equilib module of FactSage   F I G U R E 1   Images of (A) ZrB2  oxidized under AA and (B) ZrB2 under SA  conditions for different temperatures and  times and (C) ZrB2 under the SA\\xa0+\\xa0Bubbler  condition. Nominal test specimen  dimensions were 9\\xa0mm by 9-mm wide and  2.25-mm thick  NARAPARAJU et Al.    \\x0c', '|   3  7.3 using the FactPS database with the formation of all possible solid, liquid, and gas products. Calculations assumed  that the activity of B2O3 (l) was unity and that the total  pressure was maintained at 1\\xa0atm.  |   3   3.1   R E S U LT S  | Oxidation experiments  Figure\\xa01 shows the appearance of ZrB2 coupons after oxidation. Figure\\xa0 1A shows coupons oxidized under AA condition, whereas Figure\\xa0 1B shows coupons oxidized under SA  conditions. A clear difference in the oxide surfaces can be  seen between the cases. Under SA, the surfaces were covered with a glassy layer (darker regions) for all oxidation  times at 1200°C, in contrast to the brighter surfaces under  AA (except for 1h which has moderate glassy surface) conditions. At 1400°C and 1500°C, specimens exhibited a glassy  layer for the shorter oxidation periods under SA conditions.  Nevertheless, only bright oxide was present at 1400 and  1500°C in case of AA condition. Figure\\xa02 shows the SEM micrographs of the oxide surface  at 1200°C for 1h under SA and AA conditions. A few brighter  grains were embedded in a glassy pool of boria (Figure\\xa02A) in   SA, whereas areas of dense bright crystals co-existed with boria  glass for AA (Figure\\xa02B). The hole encircled in the inset could  be an evaporation route for boria or other gases produced during  oxidation. Such holes were commonly found on the scale when  examined at lower magnification (not shown). Boria is difficult  to detect using EDS, and, as a result, GDOES was used on the  darker regions of Figure\\xa0 2A with the corresponding spectrum  presented in Figure\\xa0 2C. The peaks between 249 and 250\\xa0 nm  confirm the presence of boria based on previous reports. XRD  was performed on the specimen oxidized for 4h at 1200°C  under AA (Figure\\xa02D). Monoclinic zirconia (m-ZrO2) with lattice parameters of a\\xa0=\\xa05.147\\xa0Å, b\\xa0=\\xa05.203\\xa0Å and c\\xa0=\\xa05.315\\xa0Å  was identified and indexed to JCPDS card 13-0307 from the  ICDD database. This coupon had a uniform oxide scale with no  glassy phase (ie, boria) on the surface, in contrast to the coupon  oxidized at 1200°C for 1\\xa0h EDS analysis on brighter crystals of  Figure\\xa02B confirmed the presence of zirconia (See Figure\\xa02E).  From these measurements, the darker regions were identified as  boria glass and the brighter oxide as m-ZrO2.  3.2   | Oxidation kinetics  Oxide scale thicknesses from all the tested specimens are  shown in Figure\\xa03A. The oxide scale thickness increased with   F I G U R E 2   SEM analysis of the surfaces of ZrB2 coupons oxidized at 1200°C for 1\\xa0h under (A) SA, (B) AA conditions, (C) GDOES signal  from the boria on (A), (D) XRD spectrum from the oxidized ZrB2 for 4\\xa0h at 1200°C under AA condition and (E) EDS analysis of the white crystals  shown on (B)  NARAPARAJU et Al.      \\x0c', '4   |   F I G U R E 3   (A) oxide scale thicknesses vs time for all the studied conditions, (B) (oxide scale thicknesses)2 vs time for SA condition, (C)  (oxide scale thickness)2 vs time for AA condition and (D) log(k) vs 1/T for both AA and SA conditions showing the parabolic and linear growth of  oxides of ZrB2  T A B L E 1   Measured ZrO2 oxide scale thicknesses of all the studied specimens  Oxidation  time (h)  1  4  10  Oxide scale thickness in µm  1200AA  70\\xa0±\\xa010  156\\xa0±\\xa015  225\\xa0±\\xa05  1400AA  237\\xa0±\\xa05  395\\xa0±\\xa015  564\\xa0±\\xa020  1500AA  330\\xa0±\\xa050  700\\xa0±\\xa025  984\\xa0±\\xa025  1200SA  15\\xa0±\\xa010  36\\xa0±\\xa020  41\\xa0±\\xa012  1400SA  70\\xa0±\\xa040  127\\xa0±\\xa020  454\\xa0±\\xa025  1500SA  102\\xa0±\\xa020  200\\xa0±\\xa020  871\\xa0±\\xa025  1500SA+Bubbler  331\\xa0±\\xa09  respect to time at all temperatures. Separate plots of the square  of thickness as a function of time are plotted in Figure\\xa0 3B,C  for SA and AA conditions. For SA condition, the oxidation  rate seemed to follow a parabolic trend at 1200°C (straight  line in the plot of thickness squared as a function of time),  whereas at 1400 and 1500°C, the trend appeared to be transitioning to linear after 4h. In contrast, oxidation under AA resulted in a parabolic trend at 1200°C (straight line as shown in  Figure\\xa03C) as well as at 1400°C and 1500°C with increasingly  positive curvature was observed with increasing temperatures.   Nevertheless, thinner oxide scales were observed for SA compared to AA at all the temperatures. For example at 1200°C,  zirconia scales were found to be 15\\xa0±\\xa0 10\\xa0 µm thick after 1\\xa0 h,  36\\xa0±\\xa0 20\\xa0 µm thick after 4\\xa0 h, and 41\\xa0±\\xa0 12\\xa0 µm thick after 10\\xa0 h  for SA, whereas the scales were 70\\xa0 ±\\xa0 10\\xa0 µm, 156\\xa0 ±\\xa0 15\\xa0 µm,  and 225\\xa0±\\xa05\\xa0µm thick for the same times for AA. For SA, at  both 1400°C and 1500°C, the scale thickness increased more  rapidly after 4\\xa0 h under SA conditions (from 127\\xa0±\\xa0 20\\xa0 µm to  454\\xa0±\\xa025\\xa0µm at 1400°C and from 200\\xa0±\\xa020\\xa0µm to 871\\xa0±\\xa025\\xa0µm  at 1500°C). Because of the sudden increase in the thickness   NARAPARAJU et Al.    \\x0c', 'between 4 and 10\\xa0h (Figure\\xa03B), the oxidation behavior is divided into parabolic (stage 1) and linear (stage 2) regimes. All  the measured oxide scale thicknesses for each condition are  given in Table\\xa0 1 and the corresponding rate constants have  been calculated and are given in Table\\xa02. Activation energies have been calculated by plotting the  log of rate constants as a function of reciprocal temperature.  The activation energies were 73.4\\xa0 kJ/mol for AA, 103.6\\xa0 kJ/ mol for SA stage 1, and 39.8\\xa0kJ/mol for SA stage 2.  3.3   | Oxide scale morphology  SEM cross-sectional images of the oxide scales grown after  4\\xa0 h at 1200°C and 1400°C are presented in Figure\\xa0 4, which  show noticeable differences in the oxide scale morphology  of the AA and SA conditions. The oxide scale for SA was  considerably thinner than that of AA. For AA, the scale was  156\\xa0±\\xa010\\xa0µm thick after 4\\xa0h at 1200°C and 395\\xa0±\\xa015\\xa0µm after  4\\xa0h at 1400°C. In contrast for SA, the scales were 36\\xa0±\\xa020\\xa0µm  and 127\\xa0±\\xa020\\xa0µm for SA conditions at the same temperatures  and times. Moreover, the parallel “crack” type morphology  was more apparent after oxidation in AA. Liquid boria may be  present in these “crack” type pores, which would increase the  resistance to oxygen diffusion. For oxidation in AA, the outer  ~50\\xa0µm of the scale is free of these “crack” type pores after 4\\xa0h. Figure\\xa0 5A,B show the oxide scale morphology after oxidation for 4\\xa0 h at 1500°C for AA and SA conditions. The  oxide scale was ~500\\xa0µm thicker in the case of AA than SA.  Moreover, the outer portion of the oxide scale (Figure\\xa05A,B)  contained larger pores, which might have formed due to the  evaporation of boria and subsequent sintering of zirconia  grains. The presence of parallel “crack” type pores was pronounced (about 80% of the scale) in the case of SA, whereas  only the inner part of the scale produced in AA contained such  features. The presence of columnar pores in the outer part of  the scale produced in AA suggests the probable evaporation  of boria. After 10\\xa0h, oxide scale morphologies were similar in  both conditions (see Figure\\xa05C,D) containing an outer porous  layer and an inner relatively dense layer with “crack” type  pores, although the oxide scale for AA was slightly thicker   T A B L E 2   conditions  Condition  1200 AA  1400 AA  1500 AA  1200 SA  1400 SA  1500 SA  Calculated oxidation rate constants for AA and SA   kp (µm2/h)  (Stage1:1 to 4\\xa0h)  kl (µm/h) (Stage  2: 4 to 10\\xa0h)  kp   (µm2/h)  5062  31,809  96,825  168  4032  10,099  —  55  112  |   5  than for SA (984\\xa0±\\xa0 25\\xa0 µm compared to 871\\xa0±\\xa0 25\\xa0 µm). The  outermost part of the oxide layer in SA was quite similar to  that of AA 4\\xa0h exhibiting larger columnar poles. The major difference between the two oxidation conditions other than using f lowing air in SA condition was  thought to be the presence of moisture for AA. To test this  hypothesis, a new experiment was designed to add moisture  to the f lowing SA using a water bubbler. An additional oxidation test was carried out for 1h at 1500°C and the corresponding SEM cross-sectional micrograph of the oxide layer  is presented in Figure\\xa06C. The oxide scale thicknesses of the  AA, SA, and SA\\xa0 +\\xa0 Bubbler after 1h at 1500°C are shown  in Figure\\xa0 6A-C, respectively. When water vapor was added  to SA the oxide scale thickness (331\\xa0±\\xa0 9\\xa0 µm, was shown as  + in Figure\\xa03A) matched that of ambient air (330\\xa0±\\xa050\\xa0µm).  Moreover, the surface appearance of ZrB2 coupons was also  similar as shown in Figure\\xa01A,C. The oxide scale morphology was similar for both AA and  the SA\\xa0+\\xa0Bubbler, but distinctly different from the morphology for SA. At 1400°C, the oxide scale produced in AA had  two distinct porous zones (Figure\\xa04C). The outer part of the  scale consisted of equiaxed pores (see the inset in Figure\\xa06B)  surrounding a relatively denser inner zone with “crack” type  pores nearer to the substrate. SA\\xa0 +\\xa0 Bubbler also produced  similar zones where the oxide scale close to the substrate  contained “crack” type pores. However, almost all the oxide  scale produced in SA consisted of these parallel “crack” type  pores (see Figure\\xa06A). Boria glass was likely present in these  pores before being removed by dissolution in water during  metallographic preparation of the polished cross sections.  The presence of boria in these pores would mean less evaporation of boria glass under these conditions and this implies  thinner oxide scale compared to that of AA condition. An  interesting trend in the “pore” structure with respect to oxidation time was observed for SA. After 1\\xa0 h, the scale was  predominantly “crack” type pores (after 1h). In contrast, after  4\\xa0 h the scale was composed of two distinct layers with the  outer part of the scale containing ~20% larger pores and the  inner part and about 80% “crack” type pores in the inner part  (see Figure\\xa0 5B). After 10\\xa0 h, the scale was very thick and it  contained roughly equal portions of larger pores (outer part)  and “crack” type of pores (inner part, see Figure\\xa05D).  3.4   |   Thermodynamic calculations  The volatility of boria under the conditions of the oxidation  tests was assessed using thermodynamic calculations. The  starting materials were one mole of ZrB2 solid and 2.5 moles  of oxygen gas along with the corresponding amount of nitrogen (9.4 moles assuming air was 21% oxygen and 79% nitrogen) that would be present in air. The water vapor pressure  was assumed to be zero for dry air and 0.031\\xa0atm for wet air.   NARAPARAJU et Al.      \\x0c', '6   |   (A)  (B)  F I G U R E 4   Backscattered electron  (BSE) images of cross sections of oxides  grown on ZrB2 after oxidation at (A)  1200°C AA, (B) 1200°C SA, (C) 1400°C  AA, and (D) 1400°C SA after 4\\xa0h  (C)  (D)  This value was selected because it is the pressure for saturated  water vapor pressure at 25°C,16 as would be attained by bubbling air through water at room temperature prior to entering  the furnace. For dry air, the predominant vapor species was  B2O3 (g) at all temperatures. As expected the vapor pressure  of B2O3 (g) increased with temperature from ~10−5\\xa0 atm at  1200°C to more than 10−3\\xa0 atm at 1500°C. The presence of  water vapor did not affect the predicted pressure of B2O3 (g),  but the predominant species in the presence of water was  HBO2 with a vapor pressure that was 1 or 2 orders of magnitude higher than B2O3 at all temperatures (see Table\\xa03). Previous experimental studies have shown that boria evaporates readily during oxidation of ZrB2 at 1500°C,7 indicating that  a vapor pressure on the order of 10−3\\xa0atm is sufficiently high to  allow complete evaporation under conditions typically used for  oxidation studies. The presence of water vapor promotes the formation of HBO2 with a vapor pressure above 10−3\\xa0atm at 1200°C.  Based on comparison to oxidation in dry air conditions, boria  should volatilize more rapidly in the presence of water vapor due  to the higher vapor pressure of HBO2 under these conditions.  4   |   D I S C U S S I O N  The current results agree with previous reports that indicate  that the oxidation of pure ZrB2 follows parabolic kinetics at   1200°C under SA.10,11 The kinetics become para-linear and  linear in nature with increasing temperature.4 The results  show that the oxide formed at 1200°C was a mixture of  boria (l) and zirconia (s). As temperature increased, boria  evaporation left behind a porous zirconia scale.5,6,10,11 As  long as boria is present as a continuous layer, the reaction  is controlled by diffusion through it, which we have observed in the 1200 SA condition. However, the continuous  boria (l) outer scale was not present for AA condition (see  Figure\\xa01). The loss of the protective boria layer resulted in  a loss of protection and thicker oxide scales compared to  SA. With increasing temperature, oxidation kinetics were  governed by the competition between the rates of boria formation and evaporation. The parallel “crack” type pores  that were present at lower temperatures and shorter oxidation times were presumably filled with liquid boria, which  offered protection against oxidation (Figure\\xa04). In addition,  AA has larger pores on the outer part of the coating suggests that probably higher boria evaporation than SA. At  1500°C, evaporation of boria was faster than its formation  and the presence of larger pores in the outer part of the  scale suggests that the oxidation rates have increased significantly. However, under SA, oxidation rates were lower  up to 4h and the corresponding micrographs shown in  Figure\\xa06A and Figure\\xa05B exhibit thinner oxide layers with  mostly “crack” type pores. The porous outermost oxide   NARAPARAJU et Al.    \\x0c', '(A)  (B)  (C)  (D)  |   7  F I G U R E 5   BSE cross-sectional images of cross sections of oxides grown on ZrB2 after oxidation at 1500°C for (A) AA, (B) SA after 4\\xa0h, (C)  AA, and (D) SA after 10\\xa0h. Several micrographs are stitched together to show the complete oxide scale due to their larger thickness  layer after 4\\xa0 h suggests the rapid evaporation of boria has  started and a change in the oxidation rate can be expected.  As shown in Figure\\xa0 3B, the rate change occurred between  4-10\\xa0 h. After 10\\xa0 h under SA, the oxide scale was found  to have similar pore morphology as in AA with comparable thickness. This implies that boria evaporation rate is  also influenced by the oxidation time. The larger columnar  pores in the upper part shown in Figure\\xa0 5D are similar to  that of 4h case under AA condition (Figure\\xa0 5A: that is,  faster evaporation of boria was delayed by a few hours  under the presence of synthetic air. Zhang et al used similar   experimental conditions (flowing synthetic air, at 1500°C  with a heating/cooling rates of 5°C/min) and reported that  a ~50-, 150-, and 500-µm-thick porous zirconia oxide scale  had formed after 1, 2, and 3\\xa0 h, respectively.17 In another  study, where slightly different oxidation conditions were  used (flowing synthetic air at 1500°C, 5°C/min heating,  and air quenching to room temperature) oxide scales of 30-,  60-, and 75-µm thick were reported after 1, 2 and 3\\xa0 h.18,19  Several factors influence oxidation scale thickness including air flow rate, cooling/heating rates, and measurement  sensitivity in temperature. Kuriakose et al observed that the   (A)  (B)  (C)  F I G U R E 6   BSE images of cross  sections of oxides grown on ZrB2 after  oxidation at 1500°C for (A) SA, (B) AA and  (C) SA\\xa0+\\xa0Bubbler after 1\\xa0h  NARAPARAJU et Al.      \\x0c', '8   |   T A B L E 3   Summary of predominant vapor species and vapor pressures for oxidation of ZrB2 at different temperatures  Dry Air  Predominant  species  B2O3 B2O3  B2O3  1200°C  1400°C  1500°C  With water vapor  Vapor pressure of predominant  species (atm.)  Predominant species  Vapor pressure of  predominant species (atm.)  1.3\\xa0×\\xa010−5 5.5\\xa0×\\xa010−4  2.5\\xa0×\\xa010−3  HBO2 HBO2 B2O3 HBO2 B2O3  3.5\\xa0×\\xa010−3 1.6\\xa0×\\xa010−2 5.5\\xa0×\\xa010−4 2.7\\xa0×\\xa010−2 2.5\\xa0×\\xa010−3  rate of oxidation of ZrB2 was directly proportional to the  oxygen partial pressure; however, no dependency of rate  constants on flow rates was found at 1056°C.11 The flow  rate of synthetic air has a greater influence on the oxidation  behavior at higher temperatures where the formation and  evaporation of boria played greater roles in defining the  oxidation kinetics. Our initial experiments showed a trend  of decreasing oxidation rate with increasing flow rates of  synthetic air at 1500°C, which is the subject of continuing study. In another study, the gas velocity was shown to  be an important factor in volatility, spallation, and recession of UHTC materials.13 Another important factor is the  cooling and heating rates, which affect the length of time  that specimens are exposed to oxidizing atmospheres. For  example, a 5°C/min heating rate as used in Ref. [17] added  an extra 60\\xa0mins of oxidation time in the temperature range  from 1200°C to 1500°C to the overall oxidation time of  1, 2, and 3\\xa0 h at 1500°C, compared to the higher heating  rate used in the present study. The effect was different  in another study,18 using a 5°C/min heating rate and then  quenching to room temperature at the end of the oxidation  period. As a result, huge difference in the oxide scale thickness was observed (75\\xa0µm compared to 500\\xa0µm after 3\\xa0h at  1500°C) in both the studies.17,18 To eliminate such effects,  equal heating/cooling rates of 10°C/min were used in all  the AA, SA, and SA\\xa0 +\\xa0 Bubbler conditions. As a result,  measured oxide scale thicknesses in SA condition cannot  be directly compared to other reported values without considering total oxidation time; however, our values were  similar to those reported in the studies of Zhang et al.17 Signif icant differences were seen in the oxide scale  thicknesses between AA and SA conditions at all the studied temperatures as shown in Figure\\xa0 3. The difference in  the oxide scale thicknesses can be explained by thermodynamic analysis of ZrB2 oxidation. When ZrB2 is exposed to  dry air, the predominant vapor species is B2O3 (g), whereas  other vapor species such as BO2, B2O2, BO, B2O, and B2  have signif icantly lower vapor pressures.7 The vapor pressure calculations presented in Table\\xa03 predict a shift in the  predominant vapor species and higher vapor pressures in  the presence of water vapor, which should lead to increased  evaporation of B2O3 at lower temperatures. At 1200°C, the   vapor pressure of the predominant species in the presence  of water vapor (HBO2) was about the same as the vapor  pressure of B2O3 in dry air at 1500°C. As a result of the  increased vapor pressure, measurable bor ia evaporation  star ts at much lower temperatures when water vapor is  present. The hole in the oxide scale shown in Figure\\xa02B is  evidence of such evaporation. At 1400°C and 1500°C, the difference in the vapor pressures of predominant species is smaller, but still higher in  air containing water vapor. At 1400°C, the vapor pressure  of HBO2 was about one order of magnitude higher than  B2O3. The more rapid evaporation of the bor ia led to an  increase in the thickness of the remaining zirconia scale.  The abrupt increase in the oxide scale thickness between 4  and 10\\xa0h at 1400°C suggests a transition in oxidation mechanism at longer times. Presumably, the transition occur red  when the bor ia no longer formed a continuous layer and  the underlying ZrB2 was directly exposed to the oxidizing  environment. At 1500°C in dry air, the vapor pressure of B2O3 is high  enough to enable complete evaporation in f lowing air. Since  the air is f lowing, it is not saturated in B-O species and continuous evaporation is promoted. The presence of water vapor  still promoted the formation of HBO2 which would enhance  the evaporation rate compared to SA and could be responsible for the structural differences seen in Figure\\xa0 1A. As a  result of the presence of water vapor, higher oxidation rates  with thicker zirconia scales form under static air conditions  in the entire studied temperature range. The addition of water  vapor to SA supports this hypothesis because SA containing  water vapor would promote HBO2 formation with a higher  vapor pressure than boria vapor, which would lead to the increased thickness of the zirconia layer, similar to what was  found for AA conditions.  5   |   C O N C L U S I O N S  The following conclusions can be drawn from this study:  • The presence of water vapor increases the oxidation rate of  ZrB2 by promoting volatilization of the protective B2O3 at   NARAPARAJU et Al.    \\x0c', 'lower and intermediate temperatures compared to dry air. • The increased volatilization of boria, particularly at 1400°C  and below, in the presence of water is due to the formation  of a hydrated B-O compound, specifically HBO2. • This effect is severe at 1200°C, which causes oxidation  to proceed with a linear oxidation rate in the presence of  water vapor compared to parabolic kinetics in dry air.  AC K N OW L E D G M E N T S  At Missouri University of Science and Technology, this research was supported by the Enabling Materials for Extreme  Environments  signature  areas. Alec Murchie was  supported by a Graduate Assistance in Areas of National Need  (GAANN) fellowship. The authors acknowledge Andrea  Ebach-Stahl (DLR) for her assistance in conducting GDOES  measurements. The authors acknowledge Dr Mechnich for  his assistance in water vapor experiments using bubbler.  Open access funding enabled and organized by Projekt DEAL  O RC I D  Ravisankar Naraparaju\\xa0 Gregory E. Hilmas\\xa0   https://orcid.org/0000-0002-3944-1132   https://orcid.org/0000-0002-8497-0092   R E F E R E N C E S   1. Wuchina E, Opila E, Opeka M, Fahrenholtz W, Talmy I. UHTCs:  ultra-high temperature ceramic materials for extreme environment  applications. Electrochem Soc Interface. 2007;16(4):30-6.  2. Silvestroni L, Kleebe HJ, Fahrenholtz WG, Watts  J. Superstrong materials for temperatures exceeding 2000°C. Sci Rep.  2017;7:40730.  3. Opeka MM, Talmy  IG, Wuchina EJ, Zaykoski  JA, Causey  SJ. Mechanical,  thermal, and oxidation properties of  refractory hafnium and zirconium compounds.  J Eur Ceram Soc.  1999;19(13-14):2405-14.  4. Fahrenholtz WG, Wuchina EJ, Lee WE, Zhou Y, editor. Ultra high  temperature ceramics: materials for extreme environment applications. Hoboken, NJ: JOHN WILEY & Sons, Inc.; 2014.  5. Fahrenholtz WG, Hilmas GE. Oxidation of ultra-high  temperature  transition metal diboride  ceramics.  Int Mater Rev.  2013;57(1):61-72.  6. Parthasarathy TA, Rapp RA, Opeka M, Kerans RJ. A model  for  the  oxidation  of ZrB2, HfB2  and TiB2. Acta Mater.  2007;55(17):5999-6010.  |   9   8.    7. Fahrenholtz WG. The ZrB2 volatility diagram. J Am Ceram Soc.  2005;88(12):3509-12. Jacobson N, Farmer S, Moore A, Sayir H. High-temperature oxidation of boron nitride: I, monolithic boron nitride. J Am Ceram Soc.  1999;82(2):393-8.  9. Kaufman L, Clougherty EV, Berkowitz-Mattuck BB. Oxidation  characteristics of Hafnium and zirconium diboride. Trans Metall  Soc AIME. 1967;239:458-66.  10. Tripp WC, Graham HC. Thermogravimetric study of the ZrB2  in the temperature range of 800°C to 1500°C. Solid State Sci.  1971;118:1195-9.  11. Kuriakose AK, Magrave JL. The oxidation kinetics of zirconium  diboride at high temperature. J Electrochem Soc. 1964;111:827-31.  12. Brown FHJ. Progress Report No. 20-252. Pasadena, CA: Jet  Propulsion Laboratory; 1955.  13. Nguyen QN, Opila EJ, Robinson RC. Oxidation of ultrahigh temperature ceramics in water vapor. NASA/TM-2004-212923; 2014.  14. Guérineau V, Julian-Jankowiak A. Oxidation mechanisms under  water vapour conditions of ZrB2-SiC and HfB2-SiC based materials up to 2400°C. J Eur Ceram Soc. 2018;38(2):421-32.  15. Neuman EW, Hilmas GE, Fahrenholtz WG. Processing, microstructure, and mechanical properties of large-grained zirconium  diboride ceramics. Mater Sci Eng A. 2016;670:196-204.  16. Rumble JK, editor. Handbook of chemistry and physics, 62nd edn.  CRC: Boca Raton, FL; 1981.  17. Zhang SC, Hilmas GE, Fahrenholtz WG. Improved oxidation resistance of zirconium diboride by tungsten carbide additions. J Am  Ceram Soc. 2008;91(11):3530-5.  18. Kazemzadeh Dehdashti M,  Fahrenholtz WG, Hilmas GE.  Oxidation of zirconium diboride with niobium additions. J Eur  Ceram Soc. 2013;33(10):1591-8.  19. Kazemzadeh Dehdashti M, Fahrenholtz WG, Hilmas GE. Effects  of transition metals on the oxidation behavior of ZrB2 ceramics.  Corros Sci. 2015;91:224-31.  How to cite this article: Naraparaju R, Maniya K,   Murchie A, Fahrenholtz WG, Hilmas GE. Effect of  moisture on the oxidation behavior of ZrB2. J Am  Ceram Soc. 2020;00:1-9. https://doi.org/10.1111/ jace.17500  NARAPARAJU et Al.      \\x0c']"
},{
  "_id": 47,
  "PDF": "Effect of oxidation at 1100°C on the strength of ZrB2–SiC–graphite ceramics.pdf",
  "Text": "['Journal of Alloys and Compounds 509 (2011) 6871-6875  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j a l l c o m  Effect of oxidation at 1100  C on the strength of ZrB2-SiC-graphite ceramics  Wang Zhi a , Qu Qiang b , Wu Zhanjun a,∗ , Shi Guodong a  a School of Aeronautics and Astronautics, Faculty of Vehicle Engineering and Mechanics, State Key Laboratory of Structural Analysis  for Industrial Equipment, Dalian University of Technology, Dalian 116024, China b China Academy of Launch Vehicle Technology R&D Centre, Beijing 100076, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 23 December 2010 Received in revised form 28 March 2011 Accepted 29 March 2011 Available online 5 April 2011  Keywords:  Zirconium diboride Surface oxidation Microstructure  1.  Introduction  One standard surface crack was introduced at the center of the tension surface of the test specimens with a Vickers indenter to investigate the effect of oxidation on the strength of ZrB2 -SiC-graphite ceramic. The flexural strength of the pre-cracked specimen was 371.7 MPa, which was lower than the strength of 500 MPa for the original ceramic. Oxidation in dry or moist air was employed for 30, 60, or 90 min. The flexural strength of the oxidized specimens increased as the oxidation time increased up to 60 min and then the flexural strength did not further increase. The flexural strength of specimens oxidized in dry air was greater than those specimens oxidized in moist air, which revealed that the compounds of glassy structure could better heal the cracks on the surface of the specimen than the compounds of lamellar structure. The strength of the oxidized specimen was comparable to the strength of the pre-cracked specimens.  © 2011 Elsevier B.V. All rights reserved.  Zirconium diboride (ZrB2 ) is a potential candidate for a variety of high temperature structural applications, such as furnace elements [1], plasma arc electrodes [2], hypersonic aircraft [3], reusable launch vehicles [4], or rocket engines and thermal protection structures for leading edge parts on hypersonic reentry space vehicles [5]. Because of recent efforts to develop hypersonic aerospace vehicles and re-usable atmospheric re-entry vehicles, interest in UHTCs has significantly increased in the past few years. As a result, groups in the United States, Italy, Japan, India, and China are investigating ZrB2 -based ceramics [6]. Until now, it is known that the addition of appropriate amounts of SiC particles not only enhances the mechanical properties, but also improves the oxidation resistance of ZrB2 by promoting the formation of silicate-based glasses that inhibit oxidation at temperatures between 800 and 1700  C [7]. However, unsatisfactory fracture toughness is still obstacle for them to be used widely, especially for applications in severe environments [8]. Our previous works have confirmed that the fracture toughness as low as 4.5 MPa m1/2 of the ZrB2 -SiC composites was further improved to 6.1 MPa m1/2 by adding the graphite flake [9]. The cracks on the surface of the polished specimen were found as a result of the significant pullout or desquamation of the graphite flake during the process of polishing surface because of weak bonding caused by the presence of the graphite flake within  ∗ Corresponding author. Tel.: +86 411 84706791; fax: +86 411 84706791. E-mail address: wzdlut@dlut.edu.cn (W. Zhanjun).  0925-8388/$ - see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2011.03.163  the ZrB2 -SiC-graphite ceramic. These surface slots or pits in the ZrB2 -SiC-graphite ceramic are indeed due to graphite flake pullout and/or desquamation and not due to incomplete densification [9]. Generally speaking, ZrB2 based ceramics are brittle and sensitive to cracks. As a result, the structural integrity of the ZrB2 -SiC-graphite ceramic components may be seriously affected. In order to overcome this disadvantage, there are two ways: (a) inspect carefully and repair the unacceptable cracks, (b) heal the cracks and recover strength [11,12]. For method (a), allowable cracks in ceramics are so small that it is almost impossible to detect the cracks. Moreover, structural ceramics are so brittle that repair is almost impossible. For method (b), very few studies have been made of method (b). Some ceramics containing silicon and/or aluminum element have the ability to heal a cracks, which is in favor of increases in the reliability of structural ceramic components [12]. When the ZrB2 -SiC-graphite ceramic is exposed to high temperature air, the surface of the ZrB2 -SiC-graphite ceramic begins to oxidize to oxides. For instance, ZrB2 is oxidized to ZrO2 and B2O3 above 650  C [13]; SiC is oxidized to SiO2 above 900  C [14]; Graphite flake is oxidized to CO above 500  C [15]. So, SiO2 and B2O3 glasses could be formed at high temperature. Furthermore, the SiO2 can react with B2O3 to form a borosilicate glass [16]. Crack healing behavior is very sensitive to crack-healing conditions, such as crack healing temperature and time as well as atmosphere. Therefore, in this paper, the effect of the crackhealing conditions on flexural strength of a ZrB2 -SiC-graphite ceramic was investigated in detail. The surface microstructure of a ZrB2 -SiC-graphite ceramic after oxidation at 1100  C in air was also investigated.  \\x0c', '6872  2.  Experimental  W. Zhi et al. / Journal of Alloys and Compounds 509 (2011) 6871-6875  Commercially available ZrB2 powder (2 \\u242em, >99.5%, Northwest Institute for nonferrous metal research, China), SiC (1 \\u242em, >99.5%, Weifang Kaihua Micro-powder Co., Ltd., China.) and the graphite flake (mean diameter and thickness are 15 \\u242em and 1.5 \\u242em, respectively, >99%, Qingdao Tiansheng graphite Co., Ltd., China) were used as raw powders. The powder mixture of ZrB2 plus 20 vol.% SiC plus 15 vol.% graphite flake was mixed using a planetary mill (Nanjing University planetary mill Co., Ltd., China) at 260 rotations per minute for 10 h in a polyethylene bottle using ZrO2 balls and ethanol as the grinding media. After mixing, the slurry was dried in a rotary evaporator and screened by sizing screen with 800 meshes. Then the mixture was hot-pressed at 1900   C for 1 h under a uniaxial load of 30 MPa in Ar atmosphere. The flexural strength of the specimens before and after heat treatment was tested in three point bending on 3 mm × 4 mm × 36 mm bars, using a 30 mm span and a crosshead speed of 0.5 mm min−1 . Prior to treatment, each specimen was ground and polished with diamond slurries down to a 1 \\u242em finish, and the edges of all the specimens were chamfered to minimize the effect of stress concentration due to machining cracks. A minimum number of five specimens were tested for each condition. One standard surface crack was introduced at the center of the tension surface of the test specimens with a Vickers indenter using a load of 49 N for 15 s. The polished rectangular bars were heated at the temperature of 1100   C for certain time, then slowly cooled to room temperature. The heating and cooling rates were 30   C min−1 and 10   C min−1 , respectively. Two healing environments were adopted: dry air and saturated moist air, healing time for each condition is 30, 60 and 90 min, respectively. The absolute dry air and water vapor was mixed and the relative humidities of the dry and moist air at room temperature were conditioned to 0.4% and 35%, respectively. The mixed air was pumped into the reaction chamber. After oxidation, the specimens were kept in a sealed container, which was protected from ambient moisture to prevent hydration of B2 O3 . The microstructural observations of specimen were carried out by scanning electron microscopy (SEM, FEI Sirion, Holland) along with energy dispersive spectroscopy (EDS, EDAX Inc) for chemical analysis. For the crystalline phase analysis, grazing incidence X-ray diffraction (GXRD; X’Pert MRD, Panalytical, Almelo, Netherlands) was used to determine the crystalline phases of coatings. The incidence angle for GXRD was set to 1  , which resulted in a penetration depth of less than 200 nm into the specimen.  3. Results and discussion  3.1. Microstructure  SEM images of the surface shapes and cross-sectional shapes of the crack and indentation are presented in Fig. 1. EDS and XRD analysis confirmed (not shown here) that the small darker phase was SiC and it dispersed uniformly in the lighter ZrB2 matrix, the long and narrow darker phase and the slots in the surface of the specimen were graphite. The ZrB2 -SiC-graphite ceramic showed evidence of graphite desquamation (slots) during polishing because of weak interface bonding caused by the presence of graphite in the ZrB2 -SiC-graphite ceramic [10]. It was evident from the high relative density of 99.6% that these surface pits in the ZrB2 -SiC-graphite ceramic were indeed due to graphite flake desquamation and not due to incomplete densification. As a result, the slots in the surface of specimen were detrimental to the mechanical properties and the structural integrity of the ZrB2 -SiC-graphite ceramic because ZrB2 based ceramics are brittle and sensitive to cracks on the surface of the measured specimen. The Griffith fracture criterion was applied based on the brittle fracture behavior of the ZrB2 -SiC-graphite ceramic. For sharp cracks which result in a ˛−1/2 singular stress field, the ZrB2 -SiC-graphite ceramic failure can be characterized by the Griffith fracture criterion based on fracture mechanics theory in terms of half crack length, ˛cr :  acr ≈ KIC  2  (cid:2) 2(cid:3)  (1)  where ˛cr is critical crack size for brittle fracture. (cid:2) , is the critical stress which will cause propagation of a crack-like flaw, ˛cr . KIC , is a material property referred to the fracture toughness. That is to say, the Griffith fracture criterion describes the critical flaw size that can occur without catastrophic crack propagation for brittle materials. The critical flaw size was calculated to 47.8 \\u242em based on the fracture toughness of 6.11 MPa m1/2 and flexural strength of 498.8 MPa  for the ZrB2 -SiC-graphite [7,8]. The typical crack propagation paths are inserted in Fig. 1A. The radial crack at the edge of Vickers’ indentation clearly revealed that the crack propagation path was altered by the addition of graphite flake by crack deflection, branching and crack bridging. The deformation induced by stress using Vickers’ indentation was readily observed in the cross-section as shown in Fig. 1B. The surface of the ZrB2 -SiC-graphite ceramic was oxidized to oxides at 1100  C. These main reactions during the oxidation process were expected as follows: 2C(s) + O2 → 2CO(g) ZrB2 (s) + 5/2O2 → ZrO2 (s) + B2O3 2SiC(s) + 3O2 → 2SiO2 (l) + 2CO(g) xSiO2 (l) + yB2O3 (l) → xSiO2 ·yB2O3  (3)  (5)  (2)  (4)  The micrograph of the surface of the specimen oxidized for 30 min in dry air is shown in Fig. 2A. The Vickers’ indentation on the surface of the specimen after oxidation was not found due to the formation of the oxide layers. The compounds of glassy structure were readily observed on the surface of the specimen. The EDS analysis confirmed that glass layer mainly composed of oxygen element and a small quantity of Si and Zr elements, which was consistent with the presence of B2O3 [14]. This may be due to the oxidation time limited the oxidation of SiC, indicating that 30 min was not enough for SiO2 formation at 1100  C under stagnant air because the oxidation of SiC is much slower than that of ZrB2 in this temperature regime, the SiC particles could not be obviously oxidized. The B2O3 layer is expected to contain a small quantity of SiO2 during long heating at 1100  C based on the slight oxidation of the SiC by reactions (4) and (5). Moreover, ZrO2 grains could not be detected due to a B2O3 -rich layer was found to form above the ZrO2 layer because of volume expansion upon conversion of ZrB2 to ZrO2 and B2O3 (300% volume expansion based on density calculations) and/or the mutual wetting behavior of the two materials [14]. The slots resulted from oxidation of the graphite flake were easily observed below a B2O3 -rich glass, which revealed that the slots on the surface of the specimen were not healed by a B2O3 -rich glass in 30 min due to the rapid volatilization of a B2O3 -rich glass because the B2O3 rich glass with low viscosity had a higher vapor pressure [13]. In addition, the escape of CO resulted from oxidation of the graphite flake led to that liquid B2O3 is difficult to spread to the defect. Fig. 2B shows the micrograph of the surface of the specimen oxidized for 60 min in dry air. The EDS analysis confirmed that the glass layer was composed of oxygen with small quantities of Si and Zr, which was consistent with the presence of a B2O3 -rich glass. As the oxidation time increased and B2O3 volatilized, the amount of SiC oxidation increased, which resulted in the formation of a borosilicate glass. The Zr was detected in the glass layer because of the thinner glass layer that formed due to the volatilization of B2O3 at 1100  C [13]. The presence of crossed-cracks in the glass layer was presumably ascribed to volume shrinkage during cooling. The micrograph of the surface of the specimen oxidized for 90 min in dry air was similar to that of the specimen oxidized for 60 min in dry air. Fig. 3A shows the micrograph of the surface of the specimen oxidized for 30 min in moist air. The surface of the specimen was covered with the lamellar particles and a small quantity of the glass phase which was borosilicate glass and/or unreacted B2O3 glass. As treatment time increased, the surface of the specimen did not change compared with the surface of the specimen oxidized for 30 min in moist air. The typical GXRD spectrum obtained from the surface of the specimen oxidized for 60 min in moist air is shown in Fig. 3B. Apparently, the phase analysis indicated the predominant phases for the oxidized specimen were ZrO2 , H3BO3 and a trace of  \\x0c', 'W. Zhi et al. / Journal of Alloys and Compounds 509 (2011) 6871-6875  6873  Fig. 1. SEM images of (A) surface shapes and (B) cross-sectional shapes of the crack and indentation.  Fig. 2. The micrographs of the surface of the ZrB2 -SiC-graphite ceramic after oxidation at 1100   C for (a) 30 min and (b) 60 min in dry air.  ZrB2 . The presence of a trace of ZrB2 was ascribed to thinner oxide layers because of the rapid volatilization of B2O3 at 1100  C [13]. The formation of H3BO3 of lamellar structure was due to the extreme sensitivity of B2O3 toward hydrolysis in moist air. A negative standard Gibbs free energy of reaction, B2O3 would react spontaneously with water in moist air according to Eq. (6), resulting in the formation of H3BO3 of layered structure [17]. In summary, the B2O3 glass was first formed and then it reacted with water in moist air to the H3BO3 of lamellar structure during the specimen was oxidized for  30 min in moist air.  B2O3 (s) + 3H2O → H3BO3 (s)  (cid:4)r G298 = −28.8 kJ mol  −1  (6)  1  2  After oxidation for 30 min in moist air, the specimens exposed to moist air (relative humidity of 35%) at room temperature for 24 h underwent further hydrolysis (Fig. 4). Furthermore, a small quantity of particles with lamellar structures was observed on the surface of the specimen. The borosilicate glass showed better mois Fig. 3. SEM micrograph (A) of the surface of the specimen oxidized for 30 min in moist air and the typical GXRD spectrum (B) of the specimen oxidation of 60 min in moist air.  \\x0c', '6874  W. Zhi et al. / Journal of Alloys and Compounds 509 (2011) 6871-6875  Fig. 4. The micrograph of the surface of the specimen exposed to moist air at room temperatures for 24 h.  Fig. 6. The flexural strength of the specimen.  ture resistance, lower volatility at high temperature, and higher viscosity than B2O3 glass was detected on the surface of the specimen. To evaluate the hydrolysis resistance of glass structure, the rectangular bars oxidized in moist air for 30 min were immersed in water bath at room temperature for 24 h (Fig. 5). No glassy phase was detected due to the complete hydrolysis of B2O3 . All the ZrB2 particles on the surface of the specimen were oxidized to loose ZrO2 particles and the surface of the specimen was incompletely covered with the loose ZrO2 particles. Furthermore, a small quantity of borosilicate glass was observed on the loose ZrO2 particles.  3.2.  Flexural strength  Fig. 6 shows the flexural strength of the specimen oxidized in each condition. The flexural strength of the pre-cracked specimen was 371.7 ± 29.9 MPa, which lower than the original strength of 498.8 ± 21.3 MPa. The reduction in flexural strength was ascribed to the presence of the pre-cracks because the ZrB2 -SiC-graphite ceramic is sensitive to the cracks. The flexural strength of the specimens oxidized in dry air increased from 445.9 ± 12.6 MPa to 469.9 ± 4.8 MPa as the treatment time increased from 30 min to 60 min due to the formation of the glass layer. Although cracks below the porous glass layer for the specimen oxidized in dry air for 30 min were found as a result of the oxidation of graphite flake, the strength after oxidation treatment for 30 min was still improved,  which was because the sensitivity of a ZrB2 -SiC-graphite ceramic to the cracks was passivated due to oxidation. The flexural strength of the specimens oxidized in moist air for 30 min and 60 min was, respectively, 411.1 ± 16.9 MPa and 449.3 ± 5.8 MPa, which was greater than 371.7 ± 29.9 MPa of the pre-cracked specimen. In spite of dry or moist air healing environments, the flexural strength of the oxidized specimens was recovered as the treatment time up to 60 min and then the flexural strength did not further increase, which revealed that the surface of the specimen was covered with oxide layers in the oxidation of 60 min. The flexural strength of the specimen oxidized in dry air was greater than that of the specimen oxidized in moist air, which revealed that the glassy structure was better able to heal the cracks on the surface of the specimen than the lamellar structure. The lamellar crystals consist of strongly bonded boron, oxygen, and hydrogen atoms. The atomic layers are 0.318 nm apart and held together by van der Waals forces [16]. In a sense, the lamellar crystal structure of H3BO3 is similar to the microstructure of the graphite flake, which was detrimental to the healing ability of oxide layers. Compared with 449.3 ± 5.8 MPa for the specimens oxidized in moist air for 60 min, the flexural strength for non-sealed specimens and specimens immersed in water for 24 h was further reduced to 428.2 ± 4.8 MPa and 391.7 ± 9.9 MPa, respectively, due to the hydrolyzation of B2O3 glass and the exposure of loose zirconia layer. The further decrease in strength for non-sealed specimens and specimens immersed in water for 24 h also indicated that the dry condition was more favorable than moist air condition  Fig. 5. The micrographs of the surface of the oxidized specimen immersed for 24 h in the waterbath of the room temperature, (A) and (B) were low and high magnification, respectively.  \\x0c', 'W. Zhi et al. / Journal of Alloys and Compounds 509 (2011) 6871-6875  6875  for strength recovery. Compared with 371.7 ± 29.9 MPa of the precracked specimen, the strength of the specimen oxidized in dry and moist air was recovered significantly to 469.9 ± 4.8 MPa and 449.3 ± 5.8 MPa, respectively. In order to confirm that the increase in strength of the precracked specimen after oxidation treatment was actually an effect of oxide layers, the flexural strength of the pre-cracked specimen treated in vacuum using same conditions for 90 min was greater than that of the pre-cracked specimen, whereas lower than that of the specimen treated in dry and moist airs, as shown in Fig. 6. This indicated that the increase in strength of the pre-cracked specimen after oxidation treatment was actually an effect of oxide layers. Furthermore, the flexural strength of the specimen treated in vacuum was slightly greater than that of the pre-cracked specimen immersed in water for 24 h, which indicated that the heat treatment was favorable to release the residual stress at crack tip. In addition, the original specimen (not pre-cracked) was oxidized in dry air for 90 min and the flexural strength of 519.8 ± 7.6 MPa for the oxidized original specimen was greater than original strength of 498.8 ± 21.3 MPa, which further confirmed that the increase in strength of the pre-cracked specimen after the oxidation treatment was actually an healing effect of glass layer because of the presence of the cracks on the original surface of a ZrB2 -SiC-graphite ceramic [7,18]. Based on above-mentioned results, the increase in flexural strength of the pre-cracked specimen oxidized was mainly attributed to the formation of the oxide layers. Further work on the healing behavior of oxides formed at higher temperature, such as 1200 and 1400  C is continuing to understand the healing mechanism of the different oxides.  4. Conclusions  Two oxidation conditions were used, dry air and moist air for times of 30, 60, or 90 min. The flexural strength of the pre-cracked specimens was 371 ± 30 MPa, which was lower than the strength of 499 ± 21 MPa of the original specimens. The reduction in flexural strength was ascribed to the presence of the pre-cracks. The flexural strength of the specimens oxidized in dry air increased from 445.9 ± 12.6 MPa to 469.9 ± 4.8 MPa as the treatment time increased from 30 min to 60 min due to the formation of the glass layer. In spite of dry or moist air healing environments, the flexural strength of the oxidized specimens was recovered as the treatment time up to 60 min and then the flexural strength did not further  increase, which revealed that the surface of the specimen was covered with oxide layers in the oxidation of 60 min. The flexural strength of the specimen oxidized in dry air was greater than that of the specimen oxidized in moist air, which revealed that the glassy structure was better to heal the cracks on the surface of the specimen than the lamellar structure. Compared with 371.7 ± 29.9 MPa of the pre-cracked specimen, the strength of the specimen oxidized in dry and moist air was recovered significantly to 469.9 ± 4.8 MPa and 449.3 ± 5.8 MPa, respectively.  Acknowledgments  This work was supported by the China Postdoctoral Science Foundation Funded Project (20100481220) and the Fundamental Research Funds for the Central Universities (3014-852001 and DUT10ZDG05) and the National Natural Science Foundation of China (51002019 and 91016024).  References  [1]  [2]  Fang,  J. Am. Ceram. Soc. 90 (2007)  I. Bogomol, T. Nishimura, Y. Nesterenko, O. Vasylkiv, Y. Sakka, P. Loboda, J. Alloys Compd. (2011), doi:10.1016/j.jallcom.2011.02.176. I. Bogomol, O. Vasylkiv, Y. Sakka, P. Loboda, J. Alloys Compd. 490 (2010) 557-561. [3] H.L. Wang, C.A. Wang, X.F. Yao, D.N. 1992-1997. [4] C.F. Hu, Y. Sakka, H. Tanaka, T. Nishimura, S. Grasso, J. Alloys Compd. 494 (2010) 266-270. [5] C.L. Yeh, H.J. Wang, J. Alloys Compd. 509 (2011) 3257-3261. [6] I. Bogomol, T. 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Han, Scripta Mater. 59 (2008) 1222-1225.  J. Alloys Compd. 471 (2009)  Soc.  27  Fahrenholtz, G.E. Hilmas,  J.  Eur.  Ceram.  (2007)  \\x0c']"
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  "_id": 48,
  "PDF": "Effect of oxidation on room temperature strength of ZrB2 and HfB2 based ultra high temperature ceramics.pdf",
  "Text": "['Advances in Applied Ceramics  Structural, Functional and Bioceramics  ISSN: 1743-6753 (Print) 1743-6761 (Online) Journal homepage: https://www.tandfonline.com/loi/yaac20  Effect of oxidation on room temperature strength of ZrB2and HfB2based ultra-high temperature ceramics  E. Zapata-Solvas, D. D. Jayaseelan, P. M. Brown & W. E. Lee  To cite this article: E. Zapata-Solvas, D. D. Jayaseelan, P. M. Brown & W. E. Lee (2015) Effect of oxidation on room temperature strength of ZrB2and HfB2based ultra-high temperature ceramics, Advances in Applied Ceramics, 114:8, 407-417, DOI: 10.1179/1743676115Y.0000000012  To link to this article:  https://doi.org/10.1179/1743676115Y.0000000012  Published online: 05 May 2015.  Submit your article to this journal   Article views: 220  View related articles   View Crossmark data  Citing articles: 3 View citing articles   Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=yaac20  \\x0c', 'Effect of oxidation on room temperature strength of ZrB2 and HfB2 based ultra   temperature ceramics   high  E. Zapata-Solvas*1, D. D.  Jayaseelan1, P. M. Brown2 and W. E. Lee1  The effect of oxidation on room temperature (RT) flexure strength degradation in SiC-reinforced  ultra-high temperature ceramics (UHTCs) and La O -doped UHTCs has been characterised in the  2  3  temperature range 1400-16008C for oxidation times of up to 32 h. Flaw healing was identified for  oxide scale thicknesses ,50 mm, although second phase agglomerates limit  the flaw healing  effect  to scale thicknesses of 20 mm. A linear degradation of RT strength with oxide scale  thickness was observed for oxide scale thickness .50 mm. Two oxide scale configurations have  been proposed to minimise RT strength degradation. The most promising is the scale with a  porous  layer  containing non-interconnected porosity  (85-90% dense)  of  either MeLa2O7  or  MeOxCy (MevZr or Hf).  Keywords: Oxidation kinetics, Oxidation resistance, Strength, Protective coating, Porous layer  Introduction  Ultra-high temperature ceramics (UHTCs) are those with melting points in excess of 3000uC, which enable them to  be  used  in  thermal  protection  systems  (TPSs),  high  temperature  conductors  or  refractory  applications. UHTCs include Me borides and carbides (MevZr or Hf)1. Moreover, ZrB2 and HfB2 exhibit higher thermal [*100 W (m K) conductivity at room temperature and w50 W (m K) 1900uC]2 at [v40 W (m K) respective carbides from RT to 800uC]3. Hypersonic applications, such as in sharp nose  21  (RT)  21  than  their  21  cones (SNCs) and sharp leading edges (SLEs), require a  combination of high temperature  capability  and high  temperature strength, whereas high thermal conductivity  is particularly desirable due to greater thermal transport  during exposure in high temperature  reactive  environ ments, by conduction and environment4. Furthermore,  radiation  back  to  the  the use of SNC and SLE  improves maneuverability  and maximum operating  speed  of  hypersonic  vehicles, widening  the  range  of  hypersonic and reentry trajectories. Expected service temperatures for SLE and SNC are *2000uC4.  However,  ZrB2 resistance,  and  HfB2 which  possess  relatively  poor  oxidation  could  compromise  the  structural stability of  the vehicle as they start oxidising 800uC5. Therefore, UHTC  at  temperatures  as  low as  research has focused over the last decade on improving  ZrB2 and HfB2 based composite oxidation resistance by  the following:  (i)  Adding different Si-containing compounds with  the aim of  forming solid solutions with MeO2  and incorporating different elements in the outer  protective borosilicate coating, thus increasing its  viscosity, melting temperature and oxygen diffusion coefficient6,7,8. SiC-reinforced UHTCs  are considered the most oxidation resistant and  the  baseline materials  for  the  development  of  hypersonic vehicle TPS.  (ii)  Adding different MeB2 such as TaB2, TiB2 or CrB2 as borate and silicate glasses containing  oxides of  the elements listed (group IV-VI tran sition metals), which are immiscible and lead to  phase separation, temperature9.  increasing viscosity and melt ing  In  this  case,  the  glass with  lower  oxygen  diffusion  coefficient  is  the  oxidation rate controlling phase.  (iii)  Formation  of  a  protective  refractory  coating  formed  during  oxidation  of  UHTC  as  a  consequence of  the addition of  rare earth com10.   pounds  to  SiC-reinforced  ZrB2  In  the  particular  case of La2O3 addition,  the melting  point of  the crystal phase in the dense and solid  protective coating formed (La2Zr2O7) is *2200uC, which makes it a promising candidate  for UHTCs, needing to withstand temperatures *2000uC. Moreover, pyrochlore structures have  lower  oxygen  diffusion  coefficient  than  their  respective MeO2 The development of high strength UHTCs has also  been  examined  extensively12-17,  and  RT  flexural  strengths as high as 1 GPa have been obtained for SiC reinforced UHTCs18. However, hypersonic applications  require high strengths at high temperatures, and there  have been  few  studies about  the high  temperature  mechanical behaviour of UHTCs.  11.  1Centre  for Advanced Structural Ceramics,  Imperial College  London,  London SW7 2AZ, UK 2Dstl, Porton Down, Salisbury, Wiltshire SP4 0JQ, UK  *Corresponding author, email ezapata@us.es  Ñ 2 0 1 5 I n s t i t u t e o f M a t e r i a l s , M i n e r a l s a n d M i n i n g  P u b l i s h e d b y M a n e y o n b e h a l f o f  t h e I n s t i t u t e  R e c e i v e d 9 D e c em b e r 2 01 4 ; a c c e p t e d 2 2 F e b r u a r y 2 0 1 5  DO I 1 0 . 1 1 7 9 / 17 4 3 6 7 6 1 1 5 Y . 0 0 0 0 0 0 00 1 2  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO  407  8                  \\x0c', 'Only recently have the first studies of mechanical properties at temperatures*2000uC been published19-21,  and few studies have examined the effect of oxidation on mechanical properties22-24. Tului et al.22 reported an RT  flexural  strength  reduction  between  50  and  70% of  plasma sprayed ZrB2-SiC over a graphite substrate after a 30 min exposure to air at 1800uC. Sciti et al.23 studied  the effect of long exposure times (100 h) to air from 700 to 1400uC on the RT flexural strength of ZrB2-MoSi2, which resulted in an RT flexural strength reduction of *50% (from 370 to 170 MPa) for oxidation treatments over 1200uC, while RT flexural strength was unaffected for oxidation treatments from 700 to 1200uC. Guo and Zhang24 effect of oxidation at 1500uC for  studied the  times up to 10 h on RT flexural strength of ZrB2-SiC and reported a strength increase of *120% after 0.5 h  exposure to air (from 450 to 1000 MPa). However,  this  could be related to experimental procedure, which led to (*450 MPa),  low values  of RT strength  as  a  con sequence of (i) omitting to chamfer the edges of bend test  samples leading to fracture of the sample from the edges, lowering flexural strength; (ii) using 5 mm surface finish instead of the usual 1 mm, which could be largely detri mental as strength is limited by flaw size on tensile sur faces; and (iii) impact of machining on surface flaws as,  for  example,  electrical  discharge machining  (EDM)  induces  larger  surface  flaws  on UHTCs, which  are  strength limiting, than diamond loading cutting tools15,25. Therefore, it is likely that RT flexural strength  of these materials is approximately equal to that after 0.5 h exposure at 1500uC rather than that quoted for RT  giving  such a pronounced improvement, projected to  arise from so called flaw healing where SiO2 produced  during oxidation can fill pores, voids and any other surface flaw. In a previous study, Zapata-Solvas et al.15 studied the effect of 1 h exposure to air at 1400uC on RT  flexural strength. Although a slight change on RT flex ural strength,  from 700 to 680 MPa for ZrB2-SiC comfrom 620 to 660 MPa for HfB2-SiC  posite  and  composite, was found under this condition, a change in  the defect population after oxidation was reported, with  sharp crack generation within the unoxidised volume  during oxidation—cracks that could grow upon loading. Ideally, physical properties of SNC and SLE materials  should be maintained to enable reuse of hypersonic vehicles. However, mechanical properties depend strongly  on structural surface flaws, and an understanding of  how oxidation can alter the strength of UHTC components is required to avoid undesirable accidents. The  goal of this study is to clarify which oxide layer features  play a detrimental role and how the oxidation damage  could be minimised  for UHTC  component design  purposes. Therefore, long-term exposure times of up to  to 1600uC were studied  32 h  in air  from 1400  for  different UHTCs.  Experimental  (w99%, d50*2.5 mm, r ¼ 6.085 g cm ZrB2 powder Sigma Aldrich, Gillingham, UK), HfB2 powder (w99%, d50*5.0 mm, r ¼ 10.5 g cm 23, Sigma Aldrich, Gilling(a-SiC, d50*0.7 mm, r ¼ 3.217 g ham, UK), SiC powder 99%, cm 23, Good Fellow Chemicals, Hunting(w99%, d50*10 mm, r ¼ 6.51 g don, UK) and La2O3 23, Fluka Chemicals supplied through Sigma Aldrich, cm  23,  Steinheim, Germany) were used  compositions: ZrB2 þ 20 vol.%  to  form different UHTC  ZrB2 þ 20 vol.%  SiC þ 2 wt.%  SiC  (ZS20),  (ZS20La),  HfB2 þ 20 vol.% SiC  and HfB2 þ 20 vol.% La2O3  SiC þ 2 wt.% La2O3 (HS20La). Ceramic powders were  (HS20)  processed and then sintered using spark plasma sintering (SPS), which is one of the most widely used techniques for UHTCs  sintering26, as described  in a previous  work10. The main microstructural  features of SPS  UHTCs, as shown previously15,25, are as follows: (i) high  density (w99% of theoretical relative density); (ii) ZrB2   based UHTCs contain a homogeneous dispersion of SiC throughout the ZrB   matrix  with SiC grain sizes between and 2 mm, while HfB  -based UHTCs contain an  2   *1  2  inhomogeneous dispersion of SiC throughout the HfB2 large as 20 mm; and  matrix with SiC agglomerates as  (iii) La2O3 particles are often in close proximity to SiC  particles. SPS (40 mm diameter|5 mm thick) billets were cut using EDM to obtain 25 mm|2.2 mm|1.7 mm bars;  all bar surfaces were ground using a 1200 grit diamond  platen to remove surface damage up 25 mm|2.05 mm|1.55 mm dimensions. Finally, face was polished to 1 mm, and all  to  the  tensile  edges were  chamfered with the  same  surface finish. Final dimen sions of the bars used for flexure strength characterisation were 25 mm|2.0 mm|1.5 mm. Oxidation tests  for all compositions were carried out in an open air furnace operating in isothermal mode at 1400uC for 0, 1, 4, 16 and 32 h; 1500uC for 0, 1, 4, 8 and 16 h; and 1600uC for 0, 1, 2, 3 and 4 h to study the effect of oxi dation on RT flexural strength. Heating rates of 10 K 21 and cooling rates of 20 K min 21 were set. Longer min times at 1600uC were not studied to avoid catastrophic  failure of the superkanthal heating elements. Mass gain  after oxidation was calculated by weighing before and  after oxidation. RT flexure strength was calculated in air  using a three-point bending fixture with a span length of  20 mm in a conventional deformation machine (Zwick/  Roll, Munich, Germany) operated at crosshead speed of 0.5 mm min 21 according to ASTM C 1161 standard.  Calculation and corrections of RT flexure strength are described elsewhere15. Three samples at each condition  were  tested, and mass gain and flexure  strength after  oxidation were  averaged. Cross-sections of  specimens  after flexure  testing  for  scanning  electron microscopy  (SEM)  characterisation were  prepared  using  conven tional methods  involving  successive  steps of  grinding  and polishing with diamond slurries and cloths embedded with up to 1 mm diameter particles. The specimens  were observed in an SEM equipped with a field emission  gun  (LEO 15,  JEOL, Tokyo,  Japan).  Samples were  examined in secondary electron imaging mode for oxide  scale thickness quantification and backscattered electron  image mode  for  atomic  number  contrast.  Energy  dispersive X-ray spectroscopy (EDS) was used to aid  phase identification. Fracture surfaces were analysed in  a non-contact  three-dimensional optical microscope for  surface metrology (Leica DCM8, Wetzlar, Germany).  Results and discussion  Effect of oxidation on flexure strength  Table 1 illustrates the RT flexure strengths of as fabricated,  unoxidised samples,  characterised in a previous  study,  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  408  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO 8                      \\x0c', 'Table 1  RT strength from Ref. 15  UHTC  ZS20  ZS20La  HS20  HS20La  Flexure strength (MPa)  700¡90 600¡70 620¡50 690¡40  which are within the range 600-700 MPa15. RT flexure strength versus time after oxidation at 1400uC, 1500uC and 1600uC for all UHTC compositions is shown in Fig. 1a-c (v2 h, 1500uC)  respectively. It can be seen that a short  treatment  in air produces limited degradation of flexure  strength and could even be beneficial. Figure 2 shows the  RT flexural  strength versus  the oxide  layer  thickness,  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  for HS20 in which the threshold value is *20 mm. These  values  correlate well with the previously characterised  native strength limited flaws reported elsewhere except for HS2015. Therefore, once the oxide scale thickness is  larger than the size of native strength limiting flaws, the  oxide  scale  on  its  own  contains  or  forms  the  new  strength limiting flaw. The earlier strength degradation in HS20 is related to the presence of *20 mm diameter  SiC agglomerates, and La2O3 addition delays strength degradation to higher oxide scale thicknesses (w50 mm).  A beneficial effect of oxidation has also been reported by In addition, Guo and Zhang24  others15,23,24.  suggested  that  surface flaw healing was  preservation (ZS10) and ZrB2 þ 30  after oxidation SiO -rich liquid formed during oxidation could fill pores  in ZrB2 þ 10  responsible for vol.% SiC (ZS30) ceramics, as  strength  vol.% SiC  2  revealing a linear RT flexure strength degradation after an  or microstructural damage, which is in good agreement  oxide layer threshold is reached. There are  two main  features  in flexure  strength  behaviour of UHTCs after an oxidation treatment, as  observed in Fig. 2. (i) When oxidation layer thickness is  below a threshold value, there is no an obvious effect on  RT strength, either detrimental or beneficial (oxi-dation  at 1400uC for 1 h is the only experimental con-dition at  which all compositions have slightly  improved flexure  strength). (ii) When oxidation layer thickness is above a  threshold value a linear degradation of flexure strength  occurs. The threshold value for oxide  layer thickness  corresponds    to    oxide   layer  thickness  of  *50 mm except  with this study.  Once the strength degradation threshold is reached, a  linear degradation of strength with oxide scale thickness  occurs. The greater the oxide scale thickness,  the larger  the strength degradation. Nonetheless,  the RT strength  degradation, which  is measured  as  the  slope  of  the  flexure  strength versus oxide  layer  thickness  for oxide  layer thicknesses higher than the threshold value, varies  for different UHTC compositions and temperatures as  listed  in  Table  2, which  suggests  that  oxide  scale  products may generate a new flaw population as was pointed out in a previous study15. The lines in Fig. 2 are  1  Strength versus oxidation time for all UHTC compositions at a 14008C, b 15008C and c 16008C; strength average values with  their respective error ranges are represented  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO  8  409    \\x0c', 'Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  2  Strength versus oxide scale thickness for all UHTC compositions at a 14008C, b 15008C and c 16008C; strength average  values are represented, which have error range lower than 10%; guideline that fits linear RT strength degradation has been  included  Table 2  compositions  RT strength degradation coefficients for all UHTC 14008C, 15008C and 16008C (heating rates of 10 K min 21 and cooling rates of 20 K min 21 were set during oxidation experiments)  at  Strength degradation (MPa/mm)  UHTC  ZS20  ZS20La  HS20  HS20La  14008C  15008C  16008C  0.7¡0.1 1.2¡0.2 13.6¡0.8 7.9¡0.5  1.7¡0.4 1.9¡0.2 1.5¡0.1 3.2¡0.2  3.4¡0.3 7.2¡1.1 8.2¡3.0 2.4¡0.1  a guide to  indicate data points considered for  linear  fitting of RT  strength degradation  sensitivity calculation. In addition, the data points, which were not  considered in the fitting or which are out of the guide  lines, have an oxide layer thickness below the threshold.  According to flexure strength degradation sensitivity k,  flexure strength as a function of oxide layer thickness  could be used as follows if it is considered to have a  nearly negligible effect on flexure strength in the case of  oxidation layer thickness below the threshold:  s ¼ sRT 2 kðx 2 x0Þ     x $ x0  ð1Þ  where sRT is the RT flexure strength for x,x0, x0 is the oxide layer thickness threshold and x is the oxide layer  thickness. Trends observed are as follows:  410  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO 8  (i)  SiC-reinforced UHTCs  show a lower  slope for  RT strength  degradation  (Fig.  2),  except  for  HS20, which exhibits an anomalous behaviour at 14008C and 16008C. However, behaviour to ZS20 at 15008C.  it shows a similar  (ii)  La O   2  3  -doped UHTCs have a higher  sensitivity  or  lower  damage  tolerance  for RT  strength  degradation (Fig. 2). As a result, long-term oxidation is not beneficial for  UHTC RT strength as suggested by Guo and Zhang24,  who observed an RT strength value after oxidation  much higher than RT strength after sintering in all cases  (,400 MPa for RT strength and .650 MPa in all cases  for RT strength after oxidation at 15008C). In addition,  Guo and Zhang24 observed an approximately  linear  decay for RT strength after oxidation at 15008C for  oxidation times from 0.5 to 10 h, which qualitatively  agrees well with the data presented in this study. However, the degradation observed by Guo and Zhang was  less pronounced than reported herein. The reasons for  this could be (i) the cooling rate used by Guo and Zhang  was 10 K min 21, while 20 K min 21 was used in this  study, which increases compressive stresses at the interface upon cooling and could produce larger flaws for the  same oxidation conditions;  (ii) differences  in sample  preparation  routines  that  could  lead  to a different  strength-limiting surface flaw population; and (iii) Guo  and Zhang used samples with a thickness of 2.5 mm24,  \\x0c', 'while  the samples  in  this study were 1.5 mm  thick. Therefore, although our measured oxide layer thickness at 15008C27 agrees well with their reported data, 200 mm  of oxide scale represents 27% of total thickness in our samples, while only 16% of the thickness of Guo and Zhang’s samples24 was oxidised and compressive stresses  upon cooling throughout the volume will be lower than in our samples. Strictly, to make the data comparable, it should be analysed in terms of relative values as seen in Fig. 3. Although the stress range in Guo and Zhang’s study  is not directly comparable,  it can be seen that  ZS20 stress values lie between extrapolation lines of the ZS10 and ZS30 data  from Guo and Zhang24. The  invariance of RT strength with relative sample scale means that the larger the component dimensions, the higher  the oxidation damage  tolerance of  the component, which is a promising result as real components are much  larger  than  laboratory  specimens. These results indicate that RT strength after a high temperature treatment is material composition dependent, while RT strength depends on sample preparation routine, as in our case samples were fabricated by SPS and Guo and Zhang’s  samples were  fabricated by hot pressing24.  In addition, the effect of cooling rate might be minimal as both cooling  rates are  slow and no appreciable differences are seen in Fig. 3. It is also sensible to use standard dimensions to make results directly comparable as in this study, in which samples were prepared and tested according to the ASTM C 1161 standard. Moreover, RT strength degradation sensitivity in Fig. 3 has a clear trend: the higher the SiC content in ZrB2 based UHTCs, the higher the RT strength degradation sensitivity. The presence of residual stresses in SiC-reinforced UHTCs as a consequence of  the  thermal mismatch between SiC and ZrB2 has been previously characterised by X-ray diffraction (XRD)15, Raman scattering and  neutron diffraction28. In those studies, it was observed  that the higher the SiC content, the higher the residual stresses  throughout  the UHTC bulk. These residual stresses are generated upon cooling after sintering at high temperature. The importance of stress generation on cooling and also the protective behaviour against oxidation provided by SiC addition  are well known and have  led to use of SiC contents ranging from 10 to 30 vol.%  in  baseline UHTC materials. Therefore,  residual stresses in the unoxidised UHTC bulk along or near the oxide scale  interface could have a negative impact and add  cumulatively  to  stresses generated during cooling as a consequence of thermal mismatch between the oxide scale and the unoxidised UHTC bulk. Furthermore, Guo and Zhang24 attributed RT strength  degradation to a SiC-depleted layer, which cannot be  concluded by studying just RT strength degradation at one temperature. In a previous study27, it was shown  that the SiC-depleted layer was thicker after oxidation at  14008C than after oxidation at 15008C as a consequence of boria evaporation from the SiO2 rich melt, which reduces the oxygen diffusion coefficient of the protective SiO  -rich  layer. In fact, oxidation resistance in ZS20 was poorer at 14008C  than at 15008C27, while RT  strength degradation sensitivity is smaller at 14008C than at 15008C, which indicates that absolute oxidation temperature may play an important role in the RT strength degradation and  it does not depend  solely on  the thickness of the oxide scale or particles in the oxide scale. In addition, the oxidation kinetics o  the Si O-rich melt  layer and  the porous oxide  layer are  identical and  thicknesses of    both  layers  follow an approximately constant proportion27, which indicates that RT strength  degradation could be attributed  to either  the oxide porous layer, SiC depleted layer and porous ZrO2 layer, or the whole oxide scale, in SiC-reinforced UHTCs. In fact,  Guo  and  Zhang  ,  s  data scale linearly with either the whole oxide  layer thickness as seen  in Fig. 3 or the porous oxide layer  2  f   .24  Effect of oxide scale on strength degradation  Strength  degradation  could  be  produced  by  damage  induced  by  chemical  reactions  during  the  oxidation  process  or  by  cumulative  stresses  generated  by  the  thermal mismatch of different oxidation products upon  cooling. The fact  that RT strength degradation is rela tively  size  (thickness)  independent  indicates  that RT  strength  degradation  could  be  produced  by  thermal  mismatch upon cooling. Otherwise, damage induced by  chemical reactions should be independent of oxide layer  thickness, under the same oxidation conditions, produ cing mechanical failure at similar Guo and Zhang24  stress levels, whereas  found low RT strength degradation  compared to our  results.  In addition,  the observation  that RT strength degradation sensitivity increases with  temperature  could  be  explained  in  terms  of  damage  induced either by chemical  reactions,  faster oxidation  kinetics at higher temperature and more violent energy  release from chemical reactions, or by thermal mismatch  as  stresses generated upon cooling are proportional  to  oxidation  temperature,  although  it  will  be  shown  that  thermal mismatch is  responsible  for RT strength  degradation  and  explained  in  terms  of  oxide  scale features. have  coefficients  of  thermal  MeB2  l (CTEs) of *7|10 21 for    the studied temperature range2, while  oxide  products  have CTEs  *10 21. Table  illustrates  the RT CTE  UHTC components and  some of  the  formed oxide  products. We assume there is an approximately constant  thermal mismatch Dl  (*3|10 21) between  the  oxide scale and the unoxidised UHTC bulk over the  whole temperature range studied, and the oxide scale  products  have   a   higher  thermal  expansion  coefficient   expansion                   26 K  the   of          of           25 K  3   26 K  3  Comparison of RT strength degradation from Guo and  Zhang22 and ZS20 data from this study; comparison is  made on basis of oxide scale relative to sample thickness  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  6  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO  411  8   2                                    \\x0c', 'than the unoxidised UHTC bulk. As a result,  there are  compressive stresses at  the interface upon cooling that  may produce slow crack growth as study15  indicated in a pre vious  and  subsequently RT strength  degra dation.  In addition,  stresses generated during cooling  scale linearly with oxide layer thickness if the oxide layer is homogeneous and thermal mismatch Dl is constant (at *3|10 21). Thus,  26 K  compressive  stresses at  the  interface between oxide  scale and un-oxidised volume  are, to a first approximation, proportional to oxide layer thickness, and the product of ETDl as TDl is the ther mal strain related to thermal mismatch upon cooling at  the oxide scale interface. From now on,  the discussion  will  focus on ZrB   -based UHTCs and the  impact of  2  secondary phase microstructural homogeneity, although  HfB   -based ceramics will also be mentioned.  2  SiC-reinforced UHTC oxide products  form a layer in which there is a SiO2 rich, glassy coating (i) on top of a porous layer of ZrO2 (ii) and a SiC depleted layer (iii) over the un-oxidised  structure as shown in Fig. 4a,  UHTC bulk (iv). However, La O   -doped SiC-reinforced  2  3  UHTC oxide products form a more homogeneous oxide  layer with all oxidation products (ii) and just a thin outer  coating of SiO   -rich glass with some ZrO  particles  (i)  2  2  over  the  unoxidised  bulk  (iii)  as  shown  in Fig.  4b.  Table  2  indicates  that  the higher  the oxidation tem perature, the higher the RT strength degradation. SiO  -rich glass has a thermal expansion coefficient  two orders of magnitude lower than ZrO2 particles. In  addition, a  SiO -rich glass is partially melted at the high  temperatures of  this  study. Therefore,  there  is no  contribution from  -rich glass to the compressive  SiO  forces at  the  interface of  the oxide scale. Moreover,  cracks in the  SiO  -rich glass are formed in the case of  thick coatings, and detachment of SiO   -rich  glass from  porous oxide scale is observed as seen in Fig. 5. As a consequence, RT strength degradation should be related  A   2  2  2  2  2   exclusively to the porous oxide layer, SiC depleted layer  and porous ZrO layer, in SiC-reinforced UHTCs, which has on average a thermal mismatch of *3|10  2  26 K  21.  SiC-reinforced UHTCs  are  less  sensitive  to RT  strength degradation than La O   -doped UHTCs. La O  2  3  2  3  dissolves quickly in glassy SiO2 to form La2Si2O7, as shown by Gao and Jiang31, who used La2O3 to increase creep resistance of MoSi2 containing SiO2 in grain boundaries and triple junctions. In addition, the reaction  to form La2Si2O7 was complete after SPS of MoSi2 at 1400uC for 5 min and doping concentrations as high as *5 vol.%. Furthermore,  a  2   wt.% addition led to a  complete reaction and the formation of La2Si2O7 with pyrochlore structure, while the 4 wt.% (*5 vol.%) ad dition led to the formation of two different polymorphs,  La2Si2O7  and  La2Si O5. Therefore,  a  higher  concen tration of La2O3 would not completely react, leaving unreacted La2O3 particles. However, Jayaseelan et al.10 observed the formation of La2Zr2O7 after oxidation in air at 1600uC for 1 h in 10 vol.% La O -20 vol.% SiC-  2  3  ZrB  UHTCs  from the  reaction between La O  and  2  2  3  ZrO2. In the present study, La2O3 concentration in the unoxidised bulk is lower than 3  vol.%, 1.7  vol.% for  ZS20La  and  2.8   vol.% for HS20La,  and  also  the  amount  of  SiO2 higher than residual SiO2 present at triple junctions in as  produced  during  oxidation  is much  sintered samples of MoSi2. In addition, La2O3 particles  are  located adjacent  to SiC particles, which promote  La2O3 dissolution within the SiO2 rich layer,  increasing  its  viscosity. Therefore,  no excess of La2O3  to  form  Me2La2O7 is expected at the studied concentrations and La2O3 was not detected by EDS in the range of oxidation temperature and holding time  studied.  A viscosity  increase  requires  higher  capillary  driving  forces  to reach the  top surface, which will  reduce  the  thickness  of  the  glassy  outer  coating,  but  otherwise  leading  to more homogeneous mixing between oxide  particles and glassy phase as observed in Fig. 4b and reported previously25. The general trend of RT strength  degradation sensitivity, shown in Table 2,  is for higher  RT strength degradation sensitivity at higher oxidation  temperature. In addition, La O  -doped UHTC degradation sensitivity is slightly higher at 1400uC and 1500uC 1600uC than  2  3  and  about  two  times  faster  at  SiC-re inforced UHTCs. A more homogeneous mixing between  oxide particles and SiO -rich phase could be beneficial as  2  -rich phase  could help relax compressive  internal  2  stresses between particles in the oxide scale so that RT  strength of ZS20La is higher or similar to RT strength of  4  Oxide scale of a ZS20 oxidised at 14008C for 1 h and b ZS20La oxidised at 14008C for 4 h  Table 3  RT CTE of UHTC compounds and some oxidation  products  formed  during  oxidation  at  high  temperature  Component  CTE (K21)  Oxidation  products  CTE (K21)  ZrB2  5.9610 6.3610 4.7610 8-10610  26 15  ZrO2  10610 10610 2-4610  26 29  HfB2  26 15  HfO2  26 29  SiC  26 15  SiO2  28 30  La2O3  26 15  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  412  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO 8  SiO    \\x0c', 'ZS20 for the same oxidation process duration at 1400uC 1500uC respectively. However,  and  the  oxide  scale  containing oxide particles for the same oxidation process  is  thicker  in the  case of La O   -doped UHTCs, which  2  3  negatively contributes and slightly increases the RT strength degradation sensitivity at 1400uC and 1500uC.  The RT strength degradation sensitivity of ZS20La is  double  that    of   ZS20,   which   suggests   that a new  mechanism may be active.  In a previous report about the long term oxidation kinetics27, a tendency for the oxidation kinetic exponents to change at an oxidation temperature of 1600uC was  observed as the oxide scale seemed to stabilise and the oxidation rate reduced27. This effect was attributed to  formation of new oxide products such as Si, La-Si and  metallic  oxycarbide  compounds  as  a  consequence  of  reaction with  oxygen  containing  species  and  carbon  formed during passive oxidation of SiC. The proposed  formation mechanisms for oxycarbides are still not clear  and are the subject to further investigation, but they will  involve  the  adsorption  of  environmental  O2(g)  or SiC27.  released  CO2(g)  from  passive  oxidation  of  C contents in oxycarbide particles vary from 80 to 20 at.  %, and the  spatial distribution of  the particles  is not  clear, but  it  seems  that  the C content  increases as  the  oxycarbide is deeper inside the oxide scale, which indi cates adsorption of environmental O2(g). However, some particles were found with high C content near the top  surface, which suggests that an adsorption mechanism of  released CO2(g) may  be  active  as well. Furthermore,  adsorption of oxygen containing species,  such as O2(g) released by SiC oxi from the  atmosphere or CO2(g)  dation, would increase the global density of  the oxide  scale layer. Figure 6 shows the mass gain per unit surface thickness at 1600uC. The ratio  area versus oxide layer  between the mass gain per unit  surface and the oxide  layer thickness is the mass gain per unit volume, which is  related to global density of oxide layer and the adsorp tion/volatilisation  rates  during  oxidation. Assuming  constant  adsorption/volatilisation  rates  during  oxi dation, as would be expected in SiC-reinforced UHTCs,  linear oxidation behaviour is consistent with a constant  oxide scale global density, while a deviation to positive dm/dt, where dm is the variation of mass gain per unit and dt  area  is  the variation of oxide  scale  thickness,  results  in a  global density  increase  as  is observed in  Fig.  6  for La O   2  -doped UHTCs. An increase  in the  3  global  density  of  the  oxide  scale  thickness  is  only  understood  in  terms  of mass  adsorption  or  oxygen  containing species adsorption. The adsorption trend of  La O  -doped UHTCs clear, being greater in ZrB based UHTCs than in  HfB  -based UHTCs, indicating  that the kinetics of oxycarbides particle formation are  higher  in  ZrB -based UHTCs  than  in HfB  -based UHTCs. This confirms  that  the ratio between oxide  particle volume and SiO   -rich  glass volume increases and  the relaxation of stresses throughout the porous oxide  layer  is reduced, resulting  in an  increase of  the RT  strength degradation sensitivity. Therefore, the increase  of oxide layer global density in ZS20 is detrimental to RT  strength after oxidation. Moreover, the increase of oxide  layer global density may  introduce a component of  chemical reaction damage to the RT strength degradation. On the other hand, the oxygen diffusion coefficient  for  ZrO at 1500uC is *10 and *10 for pure SiO2 32, and it  is  2  3  2  2  2  2   2   2  210 m2 s2 1  221m2 s21  could be  even  lower  for La O   -doped UHTCs as La O  is dissolved  2  3  2  3  within  SiO   -rich  phase. Therefore, mixing  of  oxide  2  products within the oxide scale in La O   -doped UHTCs  2  3  could be detrimental  in terms of oxidation resistance as  for the same oxidation temperature and time; the oxide  scale  is  thicker  due  to  the  thinner  outer  protective  coating and rapid diffusion through ZrO2.  Exceptional case of HS20La  HfO2 has lower permeability to O2(g) than ZrO2, which results in lower reactivity with oxygen of HfB2 compared 5,32. As a consequence, the SiO -rich phase in SiC  to ZrB2 reinforced HfB2 has lower B2O3 reinforced ZrB2, which results in an increase of viscosity  2  content  than in SiC  and a reduction of the O2(g) diffusion coefficient through  the  SiO   -rich phase.  In addition,  the higher  the B O  2  2  3  content  in the SiO   -rich phase,  the higher the solubility it27.  2  of ZrO2  and  other  oxides,  such  as La2O3,  in  Moreover, La2O3 and SiO2 react at temperatures as low as 1400uC, and the reaction is complete in times as short 5 min31. Therefore, La2O3 phase in HS20La is reduced and La2O3 could react with SiO2 to form either La2Si2O7 or La2Si O5. Figure 7 shows clear evidence of lower levels of SiO2 rich phase in  as  solubility  in  SiO2  rich  HS20La compared to in the oxide  scales of ZS20 or  ZS20La depicted in Figs. 4 and 5. Only small amounts of  5  Oxide scale of ZS20 oxidised at 14008C for 32 h  6  Mass gain per unit area versus oxide layer  thickness at  16008C; mass gain average values are represented, which  have an error range lower than 5%; guideline is included  to show trend of different compositions; two slope values  have been added for easy eye determination of average  density gain in oxide scale  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  2  ;   i( )    SiO   r ich layer ,    ( i i )   PPorous   ZrO   2  layer ,   ( i i i)    S iC dep leted  al  yer    and ( i    ) Un   oxid ised UHTC bu l k  v    A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO  413  8                                                    \\x0c', 'SiO  -rich phase are found on top of the oxide scale in samples oxidised at 1400uC (Fig. 7a), while only residual SiO2 pockets are detected at 1500uC and 1600uC (Fig. 7a times at 1600uC (no holding), and b). Even for short  2  there is no clear evidence of SiO   -rich phase, and initial  2  formation of SiOxCy particles was observed near the top  surface. This high reactivity between oxidation products  might be  responsible  for  the high RT strength degra1400uC and  dation 1500uC  sensitivity  value  of HS20La  at  compared  to  ZS20  and  ZS20La  values  in  Table 2. However,  the RT strength degradation sensi1600uC is  tivity  of HS20La  at  the  lowest  observed.  As described in the  ‘Effect of oxide  scale on strength  degradation’ section, RT strength degradation could be  due to chemical  reaction damage induced by chemical  reactions during oxidation or by thermal mismatch of  different oxidation components upon cooling. Figure 8a  shows the microstructure of HS20La after 2 h oxidation at 1600uC,  in which cracks near the oxide layer interface  connecting different microstructure damaged areas or  flaws at  the oxide scale  interface are  clearly observed.  In addition, Fig. 8b shows the microstructure of HS20La after 4 h oxidation at 1600uC, in which it is clear that the  crack density in the oxide scale layer has been largely  reduced  and  the  density  of HfOxCy  and La-SiOxCy  particles at  the interface of  the oxide scale and SiOxCy  particles along the oxide scale is higher than observed in  Fig. 8a. Therefore, crack healing phenomena or healing  of chemical reaction damage is responsible for the reduced RT strength degradation sensitivity at 1600uC  of HS20La.  This  behaviour  suggests  an  expansive  character of  the reaction to form oxycarbide particles.  However,  there  is  still  room for  induced damage by  thermal mismatch as the whole oxide scale layer is solid  at high temperature. Nonetheless,  the detection of oxy carbide particles with different O/C ratio indicates that it  might  be  possible  to  control C content  to minimise  damage induced by thermal  fatigue during a high tem perature  process  as  a  consequence  of  the  thermal  mismatch.  HS20 showed unstable behaviour  for RT strength  degradation,  as  a  consequence  of  the  presence  of  7  Oxide scale of a HS20La oxidised at 14008C for 4 h, b HS20La oxidised at 15008C for 4 h and c HS20La oxidised at 16008C  for 0 h  8  Oxide scale of HS20La oxidised at 16008C for a 2 h showing significant cracking and b 4 h showing few cracks  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  414  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO 8  \\x0c', 'large  (*20mm) SiC agglomerates15,27, which is miti gated by La2O3 addition. Kinetic carbides is slower in HS20La than  formation of oxy in  ZS20La  as  revealed in Fig. 6 by a higher deviation from the initial  trend, which results in lower RT strength degradation  sensitivity for HS20La.  In addition,  the formation of  oxycarbide particles stabilises the oxide scale thickness27, as revealed by comparing Fig. 8a and b in which  the oxide layer has nearly the same thickness. Stable  thickness of the oxide scale in oxidising environments,  low RT strength  degradation  sensitivity  and  tailor ability of C content of oxycarbide particles are promis ing  features  to design a novel  coating  that protects  against  oxidation,  minimising  structural  damage.  Although the oxidation mechanisms  that  involve for mation of oxycarbide particles are still not clear,  the  release of CO2(g) at  the oxide  layer  interface by the  passive oxidation of SiC to be adsorbed in the oxide  scale and form an oxycarbide particle  in addition to  O2(g) adsorption in the oxide layer  interface might be  feasible. Therefore,  it  could be possible  to design an  oxygen diffusion barrier or  trap,  in which there  is a  continuous re-adsorption of volatile species minimising  the growth of oxide scale under oxidising environments  at high temperatures.  Flaw population after oxidation  In a previous study, the formation of sharp cracks, which  acted as new strength limiting flaws after 1 h oxidation at 1400uC of the UHTCs of this study, was reported15. Sharp  cracks could produce a strength degradation if they are  larger than the surface flaws from machining. Flaws from  machining are located on the surface, and they could be  partially healed during oxidation, resulting in a higher RT  strength value if the oxidation-induced flaws are smaller  than the original native flaws or resulting in a lower RT  strength if the oxidation induced flaws are larger than the  original native flaws. Figure 9 shows optical microscopy  images of all UHTC compositions oxidation at 1400uC,  studied after  32 h  in which the flaws responsible for  fracture are always  located in the un-oxidised UHTC  bulk. In addition, the oxide layer has a flawless character,  indicating that chemical reaction induced damage does  not  take place and confirms that  the stresses generated  upon cooling in the oxidised UHTCs, as a consequence of  the thermal mismatch of different oxidation products and  the unoxidised bulk,  are  responsible  for RT strength  degradation. Therefore,  the design of  an appropriate  composition for the coating of UHTCs is important for  protecting against oxidation as well as minimising oxi dation  impact  on  structural  properties  of UHTC  9  Optical micrographs showing flaw formation during oxidation at oxide scale interface of a ZS20 oxidised at 14008C for 32 h,  b ZS20La oxidised at 14008C for 32 h, c HS20 oxidised at 14008C of 32 h and d HS20La oxidised at 14008C for 32 h; ellipses  indicate flaw location, while arrows point towards direction for crack front propagation; white dotted line indicates position  of oxide scale interface with unoxidised UHTC bulk  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO  415  8  \\x0c', 'components  and should be  taken into account when  UHTC components are  to be  reused,  for  example,  in  repeated re-entry applications.  The best schemes to minimise structural damage upon  cooling and protect against oxidation, as  illustrated in  Fig. 10, are (i) a protective layer against oxidation on  top followed by a porous  layer and (ii) a porous  layer  with  non-interconnected  porosity, which  lowers RT  strength degradation as well as acting as a protective  layer against oxidation. Candidate phases for the scheme  of Fig. 10b are Me La O or any other  rare earth pyro 2  7  chlore structure, which has low oxygen diffusion coefficient and melting point*2300-2600uC10, and MeOxCy, as they can adsorb oxygen containing species and O/C  ratio could be tailored to minimise induced damage by  thermal  fatigue.  In terms of RT strength degradation,  the use of a fully dense protective layer in contact with  bulk  unoxidised UHTCs  is  not  recommended  as  it  increases  RT  strength  degradation  sensitivity.  In addition,  the scheme shown in Fig. 10a could mini mise RT strength degradation. However,  the protective  top coating could suffer mechanical  failure after  long  exposures as a consequence of  its thickness. Therefore,  the coating should be a thin layer, which does not grow  with  long  exposure  and  has  a  low oxygen  diffusion  coefficient,  to effectively protect against oxidation and  minimise RT strength  degradation. As  a  result,  the  scheme in Fig. 10b may be a better option for providing  oxidation protection. Nonetheless, porosity has  to be  tailored as open porosity allows direct detrimental oxi dation of  the UHTC core, and the is *85-90% of  suggested density  range of  this  layer  relative density,  in  which the necks between particles have grown enough to  avoid  pore  interconnection  or  open  porosity, which  could allow O2(g)  to directly reach the UHTC core and  could  protect UHTCs  against  oxidation  acting  as  oxygen diffusion barrier or trap. Throughout this study,  it has been suggested that crack growth mainly occurs  on cooling, and Fig. 2 supports  this hypothesis as RT  degradation  scales  with  oxide  layer  thickness.  A transmission electron microscopy study is  currently  ongoing to analyse  the  exact  composition of different  MeOxCy phases of  the oxide layer and also to observe  grain boundaries of SiC-ZrB2, La2O3-SiC and La2O3- ZrB2 interfaces in the unoxidised bulk. Jayaseelan et al. showed that 6H a SiC transformed into 3 C and 15 R  SiC and hexagonal ZrB2 transforms into cubic ZrB after hot pressing34. Therefore, thermal stability of SiC-ZrB2, La2O3-SiC and La2O3-ZrB2 interfaces for long term  oxidation treatments will be analysed as they could play  an important role in physical and chemical properties at  high temperatures.  In summary, the effect of oxidation for long exposures  times on mechanical properties has been analysed and  strength degradation is characterised over the temperature range 1400-1600uC. The presence of an intermedi ate porous  layer with a protective coating on top was  found to be the combination providing minimal strength  degradation.  In addition,  it has been proposed that a  porous layer with non-interconnected porosity of either  Me La O or MeO C  on top of UHTCs could provide  2  7  x  y  protection against oxidation as well as minimising RT  strength degradation. The oxidation kinetics of MeOxCy  phases need investigation to better understand how they  can  be  used  for  the  benefit  of UHTCs. Results  are  obtained in a mild temperature range for UHTC oper ating temperatures and cannot be directly extrapolated  to hypersonic conditions. However, insights on induced damage after oxidation at  the  test conditions could be valuable for design purposes of UHTC components.  Conclusions  The effect of oxidation on RT strength has been studied,  and  RT  strength  degradation  was  characterised.  Although  3O -doped UHTCs reveal more protective  La2  behaviour against oxidation at 1600uC, the RT flexural  strength could be significantly reduced, as a consequence  of the average density increase of oxide scale, depending  on where oxycarbide particles are  formed. A  linear  degra-dation for RT strength was found for oxide layers  thicker than 50 mm. In addition, the presence of large SiC  agglomerates in HS20 increased strength degra-dation.  The higher the oxidation temperature, the higher the RT  strength degradation sensitivity k as a con-sequence of  thermal mismatch between oxide layer and unoxidised  volume. In addition, the longer the exposure time at the  same oxidation temperature, the lower the retained RT  strength. Two  configurations minimise RT  strength  degradation:  10  Schemes to minimise RT strength degradation: a protective bilayer scheme; b protective single layer scheme  Zapa ta -So lvas e t a l .  E f fec t o f ox ida t ion on RT s t reng th o f UHTCs  2  2  416  A d v a n c e s i n A p p l i e d C e r am i c s  2 0 1 5  VO L 114  NO 8  \\x0c', '(i)  The presence of a porous intermediate layer with  a  thin protective  coating  against oxidation on  top. The formation of MeOxCy (MevZr or Hf) at the interface between oxide scale and unoxidised  (ii)  volume, which  also mitigates  the  undesirable  effect of SiC agglomeration.  In  conclusion,  a  protective  layer with  non-inter connected  porosity,  such  as Me La O  and MeO C ,  2  7  x  y  could minimise degradation of mechanical properties.  Me La O phases are already being used in turbines and  2  7  jet engines to protect metal alloys from oxidation, and  MeOxCy phases could act as oxygen containing species  absorbers and diffusion barriers, providing a promising  new approach to protect a system from oxidation at high  temperatures. Furthermore, O/C ratio could be tailored  to produce an oxycarbide with the  same CTE as  the  UHTC matrix, minimising RT strength  degradation  upon cooling.  Acknowledgements  We  acknowledge  Professor Mike Reece  (Nanoforce  Technology Ltd, Queen Mary, University of London,  UK) for providing access to the Spark Plasma Sintering  facility. E.Z.S. acknowledges the support of ‘Fundacio´ n  Ramo´ n Areces, Spain’  and the Centre  for Advanced  Structural Ceramics  (CASC)  for his postdoctoral  fel lowship to stay at Imperial College London, UK. D.D.J.  acknowledges the support of DSTL, UK, for providing  the  financial  support  for  this work  under  contract  number DSTLX-1000015413. 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},{
  "_id": 49,
  "PDF": "Effect of pre-oxidation on the ablation resistance of ZrB2–SiC coating for SiC-coated carbon-carbon composites.pdf",
  "Text": "['Available online at www.sciencedirect.com  Ceramics International 41 (2015) 2582-2589  CERAMICS  INTERNATIONAL  www.elsevier.com/locate/ceramint  Effect of pre-oxidation on the ablation resistance of ZrB2-SiC coating for SiC-coated carbon/carbon composites  YuLei Zhangn, Zhixiong Hu, Boxing Yang, Jincui Ren, Hejun Li  State Key Laboratory of Solidification Processing, C/C Composites Research Center, Northwestern Polytechnical University, Xi’an 710072, PR China  Received 14 October 2014;  received in revised form 2 November 2014; accepted 3 November 2014  Available online 11 November 2014  Abstract  To improve ablation resistance of carbon/carbon (C/C) composites at high temperature, ZrB2-SiC coating was prepared on surface of SiCcoated C/C composites by supersonic atmosphere plasma spray. The coated C/C composites were pre-oxidized in air at 1373 K. In the present work, the influence of pre-oxidation treatment on the microstructure and ablation resistance of ZrB2-SiC coating was investigated. Ablation resistance of coated C/C composites was tested in oxyacetylene torch environment with the heat flux of 2400 kW/m2. Results show that the liquid glass phase seals defects in the as-sprayed coating after pre-oxidation for 15 min and a dense coating is obtained. The pre-oxidation coating can protect C/C composites for 90 s, which is due to the formation of a dense oxides layer and the molten ZrO2 layer. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  Keywords: Carbon/carbon composites; Coating; Supersonic atmosphere plasma spraying; Pre-oxidation; Ablation  1.  Introduction  Carbon/carbon (C/C) composites are promising candidates for aerospace application (nose tips, leading edges, and nozzles of solid rocket motors) as high-temperature thermal field the oxidation of C/C composites materials [1-3]. However, above 723 K in air cannot meet the demand of practical application requirement, especially for high-temperature (higher than 2200 K) associated with high-speed combustion gas flow in oxidizing environment [4-6]. Hence, more efforts need to be made to improve the ablation resistance of C/C composites in such extreme environments. It is well known that coating is an effective solution to protect C/ C composites at high temperature. In recent years, many ultra-high temperature ceramics (UHTC), such as ZrC, HfC, and HfB2, have been used as anti-ablation coatings for C/C composites [7-10]. As one of the UHTC, ZrB2 has high melting point (3313 K), high thermal conductivity (65-135 W/m K), relative low density (6090 kg/m3) and excellent thermal shock resistance, which is  nCorresponding author. Tel.: þ 86 29 88491384; fax: þ 86 29 88492624. E-mail address: Zhangyulei@nwpu.edu.cn (Y. Zhang).  http://dx.doi.org/10.1016/j.ceramint.2014.11.006  0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  suitable for aerospace application [11-13]. However, it is not feasible for ZrB2 as anti-ablation coating due to its rapid oxidation at high temperature [14,15]. In previous studies, ZrB2 has been modified by SiC, where ZrB2 provides mechanical resistance and SiC generates an oxidation protective scale on the material surface at high temperature [16-21]. In addition, ZrB2 with an addition of 20-30 vol% SiC exhibits good property at high temperature [19,21,22]. Supersonic atmosphere plasma spraying (SAPS) has been applied to prepare the UHTC coatings [19,22]. The temperature of plasma arc is about 10,000 K and velocity of particle is up to 600 m/s. The deposition of melted particles on the substrate can generate great jet impact force, which is beneficial to the formation of dense coating with good bonding strength [23]. Yao et al. [24] prepared the ZrB2-SiC coating by SAPS, which showed good ablation resistance after ablation for 60 s. While the microholes and microcracks on surface of the assprayed coating limited it for further ablation test. Previous results implied that thermal pre-treatment of ZrB2-SiC in air at temperature between 1373 and 1473 K could be beneficial for its oxidation resistance [25]. Tului et al. [22] reported that ZrB2-SiC coating had better property after pre-oxidation in air  \\x0c', 'Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  2583  at 1373 K, which was attributed to formation of self-protection scale in oxidizing environment. In present work, the ZrB2-SiC coating was prepared by SAPS. The SiC inner layer was first prepared by pack cementation on surface of C/C composites to relive the mismatch of the coefficient of thermal expansion (CTE) between ZrB2-SiC coating and C/C composites. The coated composites were pre-oxidized in air at 1373 K. The purpose of this study was to describe the influence of pre-oxidation on the microstructure and ablation resistance of the coating. The ablation mechanism of the coating under oxyacetylene torch environment was also discussed.  2. Experimental  2.1. Preparation of ZrB2-SiC coating  The C/C composites were prepared by Thermal gradient chemical vapor infiltration (TCVI) with a density of about (Ф30 \\x02 10 mm) were 1700 kg/m3, and the samples cut from bulk 2D C/C composites. Then the samples were cleaned in an ultrasonic bath with ethanol and dried at 373-423 K for 2 h. The SiC inner layer was prepared by pack cementation with Si, C and Al2O3 powders in inert atmosphere at 1973-2173 K for 2 h [26]. ZrB2-SiC coating was prepared on the SiC-coated C/C composites by SAPS in air. The ZrB2 and SiC particles were selected as the raw material. Commercially available ZrB2 powders (purity 4 99.9%, 800 mesh) was supplied by DanDong Research Insitute of Chemical Industry. SiC powder (purity 4 99.9%, 800 mesh) was supplied by Linyi Jin meng Carborundum Co. Ltd. The mixture particles containing 80 vol % ZrB2 and 20 vol% SiC were preliminary prepared by attrition milling with ZrO2 milling media for 2 h. However, the mixture powders cannot be used as spraying powders, because they are lack of flowability when transported from powder feeder to nozzle. In order to increase flowability of particles, spraying powders were obtained by agglomerating a kind of slurry through a spray drier. This slurry was composed of distilled water (49 wt%), polymeric binder (2 wt%), and mixture particles (49 wt%). The polymeric binder was used as an additive to agglomerate the powders. Fig. 1 shows SEM  Fig. 1. SEM image of agglomerated powders.  Table 1  Details of  the spraying parameters for ZrB2-SiC multilayer coating.  Content  Spraying current  (A)  Spraying voltage (V)  Primary gas Ar  (L/min)  Carrier gas Ar  (L/min)  Second gas H2 (L/min) Powder feed rate (g/min)  Spraying distance (mm)  Injector  internal diameter  (mm)  Injector position  Parameters  380-415  130-150  75  10  5  20  100  5.5  Perpendicular  to samples  morphology of the agglomerated powders. The average size of the powders was about 50 μm. The spraying system consists of plasma torch, powder feeder, gas supply system, water-cooling circulator, control unit with PC interface and power supply unit. The samples were perpendicular to plasma arc and injector, and the inner diameter of injector was 5.5 mm. The argon (Ar) was used as both the primary gas and carrier gas, and the hydrogen (H2) as the secondary gas. More details of spraying parameters are summarized in Table 1.  2.2. Pre-oxidation and ablation test  The pre-oxidation test concentrated on exposing coated C/C composites in ambient air at 1373 K in an electrical furnace. The as-sprayed composites were cleaned in an ultrasonic bath with ethanol and dried at 373-423 K for 2 h, and then put into isothermal region of the furnace. The oxidation time was 10, 15 and 20 min. The ablation resistance of the pre-treatment samples was carried out in oxyacetylene torch system according to the National Standard of Ablation Test method of ablative materials (GJB 323A-96) with heat flux of 2400 kW/m2 [27]. The pressure and flux of oxygen were 0.4 MPa and 0.244 L/s, and they were 0.095 MPa and 0.167 L/s for acetylene. The inner diameter of the oxyacetylene gun tip was 2 mm and the distance between the gun tip and the sample was 10 mm. In this study, all of the as-sprayed and pre-treatment samples were tested in the same condition. The ablation gun was primarily ignited. As the flame was stable, the sample was vertically placed to the flame and exposed to the flame for 60 s, 90 s and 120 s, respectively. The surface temperature of the sample was detected by an infrared thermometer.  2.3. Characterization  The crystalline structure of the coating before and after ablation was measured by X-ray diffraction (XRD, X’Pert Pro MPD) with a Cu Kα radiation of wavelength 0.154 nm. The morphology and chemical composition of the coated composites were investigated by scanning electron microscopy (SEM, JSM 6460) equipped with energy dispersion spectroscopy (EDS).  \\x0c', '2584  Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  defects. So the composites which are pre-oxidized for 15 min are selected for further research. According to the XRD pattern (Fig. 4), the coating is mainly covered with ZrO2, and the intensity of ZrB2 and SiC peak significantly decreases after pre-oxidation for 15 min. Fig. 5 shows backscatter cross-section image and EDS line analysis of C/C composites with the pre-oxidation coating. From the EDS line analysis, the content of oxygen starts to increase at the end of the coating, which implies that the oxides layer appears above the ZrB2-SiC layer. So we divide the multilayer coating into three layers: the oxides layer, the ZrB2-SiC layer prepared by SAPS and SiC inner layer prepared by pack cementation. Because the thickness of oxides layer is thin, so it is not marked in Fig. 5. The thickness of the outer layer and SiC layer is about 140 and 20 μm, respectively. In addition, no obvious crack or void is found at the interface between inner layer and outer layer, indicating a good interaction between them.  Fig.  2. Comparison  between XRD patterns  of  agglomerated  powders  and  sprayed coating.  3. Results and discussion  3.1. Phase composition and microstructure of  the coating  3.2. Ablation resistance of  the coatings  3.1.1. Phase composition of  the sprayed coating  The comparison XRD results of spray dried ZrB2-SiC powders and as-sprayed coating are shown in Fig. 2. It is evident that the intensity of SiC peak decreases and ZrO2 phase appears, which are most likely due to the oxidation of SiC and ZrB2 in air at high temperature during the spraying process.  3.1.2. Microstructure  of  the  sprayed  and  pre-oxidation  coating  Fig. 3 shows the surface SEM images of the as-sprayed and pre-oxidation coating. As shown in Fig. 3a, small voids distribute on surface of the sprayed ZrB2-SiC coating. During the spraying process, the oxidation of SiC and ZrB2 produced gaseous byproducts (such as SiO and B2O3) at high temperature, and evaporation of gases might lead to the formation of small voids. These voids provide the channel through which oxygen can diffuse to the underlying coating or C/C matrix [8]. Fig. 3b-d shows the surface SEM images of ZrB2-SiC coating after pre-oxidation for 10, 15 and 20 min. With the increase of pre-oxidation time, the content of liquid phase first increases and then decreases. After pre-oxidation for 10 min, the voids significantly decrease on surface of the coating (Fig. 3b). However, some voids are also visible because the amount of oxides is not sufficient to seal all of the voids. After preoxidation for 15 min, the voids have disappeared completely (Fig. 3c). The as-formed coating is dense and covered with a great deal of liquid phases and solid particle phases. Oxidation of ZrB2 and SiC can produce amorphous B2O3, SiO2 and crystallized ZrO2 [28-31]. The liquid phase consists of B2O3 and SiO2, and the solid phase is predominant composed of ZrO2. The dense oxides layer can reduce inward diffusion of oxygen. But the microcracks and voids appear again after preoxidation for 20 min (Fig. 3d). Long-term pre-oxidation might result in the evaporation of B2O3 and appearance of some  To study influence of pre-oxidation on the ablation resistance of ZrB2-SiC coating, both the as-sprayed and pretreatment composites were tested in oxyacetylene torch environment with the heat flux of 2400 kW/m2. The highest temperature was 2723 K. Fig. 6 shows the macrographs of the coated samples after ablation for different time. As shown clearly, all of the ablated samples are covered with white phase in the ablation center. The cracking appears on surface of the as-sprayed coating after ablation for 60 s, which suggests that the as-sprayed coating cannot provide long-term ablation resistance for C/C composites (Fig. 6a). From Fig. 6b and c, it is apparently that the pre-treatment coating is unbroken after ablation for 60 and 90 s. However, cracking also happens in the ablation center after ablation for 120 s.  3.3. Ablation morphology of pre-oxidation coating  XRD patterns of the pre-oxidation coating ablated for different time are shown in Fig. 7. As shown clearly, the XRD patterns are similar to each other, and the coating is mainly composed of ZrO2. Because the temperature in brim region is lower than that in center region, SiC is not completely oxidized and also exists in this region. Although there is no clearly difference in phase composition, the microstructure in ablation center is different. Fig. 8 shows the SEM images and EDS analyses of ablation center after ablation for different time. From Fig. 8a, the center region is covered with molten ZrO2. Due to the solution of SiO2 into ZrO2, the melting point of ZrO2 decreases, thus leading to the formation of molten ZrO2 [14,18,24]. What’s more, this molten ZrO2 with low thermal conductivity can act as thermal barrier layer. However, some bubbles can be seen on surface of the coating. During ablation process, some gases (B2O3, SiO) produce. With the increase of ablation time, the vapor pressure of the gases increases, and the escape of these gases results in the formation of bubbles [28,32]. After ablation for 90 s, the  \\x0c', 'Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  2585  Fig. 3. Secondary electron micrographs and EDS analyses of coated C/C composites:  (a) sprayed coating;  (b),  (c) and (d) pre-oxidation at 1373 K for 10 min,  15 min and 20 min,  respectively; spot energy dispersive spectroscopy analyses of  (c).  attribute to the volume expansion induced by conversion of ZrB2 to ZrO2. Although the cracking happens after ablation for 120 s, it is necessary to show the structure of the ablation center. Compared with Fig. 8b, evolution of microstructure is significant in Fig. 8c, the bubbles and cracks are obvious. In addition, the long-term mechanical denudation and the volume expansion result in the formation of coarse cracks during ablation process. In order to analyze of the failure reason of the coating after ablation for 120 s, the outer molten ZrO2 layer in center region is removed. Fig. 9 shows the SEM images and EDS analysis of the bulk. As shown clearly, the bulk is covered with dark phase, white phase and holes. By EDS analysis, the dark and white phases are distinguished as SiO2 and ZrO2, respectively. The carbon fiber can be observed clearly confirming that the coating cannot protect C/C composites any more. The cross-section backscatter micrographs of the preoxidation coated C/C composites after ablation for different time are present in Fig. 10. It can be seen that thickness of the coating has no significant change before and after ablation.  Fig. 4. Surface XRD pattern of coated C/C composites after pre-oxidation at  1373 K for 15 min.  diameter of the bubbles becomes bigger and cracks are visible in Fig. 8b, which indicates that the coating suffers great shearing force during ablation process. These microcracks may  \\x0c', '2586  Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  Fig. 5. Backscattering electron micrograph of cross-section of coated C/C composites after pre-oxidation at 1327 K for 15 min and EDS line analysis:  (a) BSE  image;  (b) EDS line analysis.  Fig. 6. Morphology of coated C/C composites after ablation: (a) as sprayed composites ablation for 60 s; (b), (c) and (d) pre-oxidation composites ablation for 60 s,  90 s, and 120 s,  respectively.  After ablation for 60 s, the coating consists of ZrO2 layer, unaffected ZrB2-SiC layer and SiC inner layer. This is not different from the previous three or four distinct layers [28,30], because the temperature in this study is much higher [18]. Combined with the EDS line analysis, the thickness of ZrO2 layer and unaffected layer is about 50 and 70 μm after ablation 60 s (Fig. 10a and b). In addition, the voids and cracks are along with this ZrO2 layer. Apparently, the thickness of the ZrO2 layer increases after ablation for 90 s (Fig. 10c). The voids have grown up and the crack propagation has taken place, which presumable results from the mechanical denudation and volume expansion derived from the conversion from ZrB2 to ZrO2. As shown in Fig. 10d, the outmost layer has debonded from C/C matrix, and SiC inner coating has  worn out after ablation for 120 s, which is in agreement with the bulk surface analysis shown in Fig. 9.  3.4. Ablation mechanism of  the coated C/C composites  There are mainly two kinds of ablation mechanism about C/C composites under oxyacetylene torch environment: chemical erosion and mechanical denudation [9]. According to the above analyses, we put forward the ablation schematic diagram of pre-oxidation coated C/C composites, as shown in Fig. 11. After pre-oxidation, the C/C composites are covered with three layers, i.e. protective oxides layer, ZrB2-SiC deposited layer and SiC inner layer (Fig. 11a). The highest temperature on the coating surface is 2723 K. In the initial stage of ablation test,  \\x0c', 'Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  2587  of cracks (Fig. 11c). As the ablation goes on, the cracks spread, and oxygen penetrates into coating through these voids and cracks, thus resulting in further oxidation of coating. As the  the protective oxides can reduce oxygen diffusion into inner layer and evaporation of the B2O3 consumes much heat, which is beneficial to ablation resistance of pre-oxidation coating. As the oxides are consumed out, the ZrB2 and SiC underlying oxides layer are oxidized further. After ablation for 60 s, the ablation center is covered with dense molten ZrO2 layer (Fig. 8b). Meanwhile, the active oxidation of SiC occurs to form SiO at high temperature. The evaporation of SiO and B2O3 leads to the formation of bubbles, which provides diffusing tunnel for oxygen. What’s more, the viscosity of ZrO2 decreases at high temperature, as a result, the bubbles in center region grow up under the shearing force of gas flow (Fig. 11b). In addition, the volume expansion derived from the conversion from ZrB2 to ZrO2 may contribute to the formation  Fig. 7. Comparison between XRD patterns of  the pre-oxidation coated C/C  composites ablation for different  time (60 s, 90 s, 120 s).  analyses:  analyses.  (a) BSE image;  (b)  high magnification  of  (a);  (c)  and  (d) EDS  Fig.  9. Backscattering  electron micrographs  of  pre-oxidation  composites  ablated for 120 s  in center  region after  removal of outmost  layer  and EDS  Fig. 8. Secondary electron micrographs of ablation center of pre-oxidation coated C/C composites and EDS analyses: (a) ablation for 60 s; (b) ablation for 90 s; (c)  ablation for 120 s;  (d) EDS analyses.  \\x0c', '2588  Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  Fig. 10. Backscattering electron micrograph of cross-section of pre-oxidation coated C/C composites after ablation for different time and EDS line analysis: (a) BSE  image of ablation for 60 s;  (b) EDS line analysis of  (a);  (c) and (d) BSE images of ablation for 90 s, 120 s,  respectively.  Fig. 11. Schematic of  the ablation process of pre-oxidation coated C/C composites.  cracks extend to the inner SiC, the active oxidation of inner SiC coating is serious, as a consequence, the thickness of inner coating decreases. In addition, the evaporation of SiO from the  inner layer to the outmost layer may also accelerate the mechanical denudation of the coating under this highpressure gas flow. After ablation for 120 s, the SiC layer is  \\x0c', 'consumed out in ablation center, only a small amount of SiO2 leaves on the surface of C/C matrix, which is confirmed in Fig. 10. Because of the mechanical denudation and mismatch (10.8 \\x02 10 \\x00 6 K \\x00 1) of CTE between the ZrO2 (1-2 \\x02 10 [18] and C/C \\x00 6 K \\x00 1) matrix [1], debonding is easy to occur between the ZrO2 layer and C/C matrix (Fig. 11d), as a result, the pre-oxidation coating is failure after ablation for 120 s. The further research about how to improve the long-term ablation resistance of the coating is needed.  4. Conclusions  ZrB2-SiC coating was prepared by SAPS on the surface of SiC-coated C/C composites. A dense coating is obtained after pre-oxidation for 15 min in air at 1373 K. The pre-oxidation coating can protect C/C composites for 90 s and then fails at 2400 kW/m2. Good 120 s under the heat flux of ablation resistance of the pre-oxidation coating is most likely to contribute to two factors: one is that the dense oxides layer reduces inward diffusion of oxygen in initial stage of the ablation test, the other is that the molten ZrO2 acts as thermal barrier layer in the next ablation test. However, the mechanical denudation and mismatch of CTE should be responsible for the failure of the pre-oxidation coating after ablation for 120 s.  Acknowledgements  This work has been supported by the National Natural Science Foundation of China under Grant nos. 51272213, 51221001 and 51202093, and supported by the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant no. 98-QZ-2014).  References  [1] T.L. Dhami, O.P. Ball, B.R. Awasthy, Oxidation-resistant carbon-carbon composites up to 1700 1C, Carbon 33 (1995) 479-490.  [2] Y.J. Wang, H.J. L, Q.G. Fu, H. Wu, L. Liu, C. 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Halloran, Oxidation of ZrB2-SiC: content on solid and liquid oxide phase formation, J. Am. Ceram. Soc. 92  influence of SiC  (2009) 481-486.  [32] W.G.  Fahrenholtz,  Thermodynamic  analysis  of  ZrB2-SiC oxidation: J. Am. Ceram. Soc. 90 (2007)  formation  of  a  SiC-depleted  region,  143-148.  Y. Zhang et al.  / Ceramics International 41 (2015) 2582-2589  2589  \\x0c']"
},{
  "_id": 50,
  "PDF": "Effect of SiC Addition on Oxidation Behavior of ZrB2 at 1273 K and 1473 K.pdf",
  "Text": "['Oxid Met (2016) 85:311-320  DOI 10.1007/s11085-015-9585-9  O R I G I N A L P A P E R  Effect of SiC Addition on Oxidation Behavior of ZrB2 at 1273 K and 1473 K  Lihua Zhang1  • Kazuya Kurokawa2  Received: 28 June 2015 / Published online: 23 July 2015 Ó Springer Science+Business Media New York 2015  Abstract  The oxidation behavior of ZrB2-SiC composites with different contents  of SiC addition was investigated at 1273 and 1473 K in air  for 12 h in this study.  The SiC addition contents ranged from 0 to 30 wt%. The results showed that when  ZrB2-SiC composites were oxidized at 1273 K in air, a two-oxide layer-structure  forms: a continuous glassy layer and a ZrO2 layer contained unoxidized SiC. When  SiC content  is 5 and 10 wt%,  the glassy layer is mainly composed by B2O3. When  SiC content is 20 and 30 wt%, a borosilicate glass could be formed on the top layer,  which could improve the oxidation resistance of ZrB2. When ZrB2-SiC composites  were oxidized at 1473 K in air, the oxide layer was composed of ZrO2 and SiO2 and  unreacted SiC. Additionally, when SiC addition content was higher than 10 wt%, a  continuous borosilicate glass layer could be formed on the top of the oxide layer at  1473 K. With  the  increase  of SiC content  in ZrB2,  the  oxide  layer  thickness  decreased at both 1273 and 1473 K.  Keywords  Zirconium diboride \\x01 Silicon carbide \\x01 Oxidation \\x01 High temperature  & Lihua Zhang  zhanglihua@eng.hokudai.ac.jp  Kazuya Kurokawa  kurokawa@office.tomakomai-ct.ac.jp  1  2  Center for Advanced Research of Energy Conversion Materials, Hokkaido University, Kita 13  Nishi 8, Kita-ku, Sapporo 060-8628, Japan  Tomakomai National College of Technology, 443 Nishikioka, Tomakomai 059-1275, Japan  123  \\x0c', 'Introduction  Ultra-high temperature ceramics  (UHTCs) have been proposed as candidates  for  applications such as thermal protection systems for hypersonic aerospace vehicles  [1, 2]. Among the UHTCs, zirconium diboride has the lowest theoretical density (6.09 g cm-3), and has good thermal shock resistance because of (65-135 W m-1 K-1)  its high thermal  conductivity  [3]. These  attributes make  it  attractive  for  aerospace applications such as thermal protection materials on hypersonic aerospace  vehicles  and  re-usable  atmospheric  re-entry  vehicles. However, the oxidation 1400 °C due  resistance  of  ZrB2  is  very  poor  at  temperatures  above  to  the  volatilization of B2O3, which results in formation of a porous, non-protective ZrO2  layer  [4, 5]. Numerous investigations to improve the oxidation resistance of ZrB2  have been reported [4, 6-11].  It was  found that SiC addition provided improved  oxidation  resistance  by  promoting  the  formation  of  borosilicate  glass.  This  borosilicate glass afforded more oxidation protection than B2O3  since it  is more  viscous, has higher melting temperature  and lower vapor pressure,  and is more  protective against oxygen diffusion [12, 13].  The oxidation behaviors of ZrB2-SiC composites in air have been well-defined  previously [6, 8-10, 14-16]. When ZrB2-SiC composite is exposed to oxidizing environments, SiC phase begins oxidizing appreciably between 1100 and 1300 °C,  resulting  in  the  formation  of  a  silica-rich  glassy  layer  that  imparts  significant  improvement to the oxidation resistance of the diboride at higher temperatures. The 1500 °C as  SiO2-containing  scale  remains  protective  up  to  at  least  SiO2  is  significantly less volatile than B2O3 at  these temperatures. Thus, ZrB2-SiC exhibits  passive oxidation behavior with parabolic kinetics over a much greater temperature  range than has been reported for pure ZrB2. Rezaie et al.  [17] also reported the  oxidation behavior of ZrB2 with 30 vol% SiC addition at  low partial pressure of  oxygen. However,  the SiC content discussed in these studies is usually in the range  of 20-30 vol%, and the influence of different SiC addition on the oxidation behavior  of ZrB2 is rarely reported [18]. Therefore,  the oxidation of ZrB2 with different SiC  contents is studied in this paper.  Experimental Procedures  Commercially available ZrB2 and SiC powders were used as  raw materials. The  ZrB2-SiC powders with different SiC content (0, 5, 10, 20, and 30 wt%) and 1 wt%  B4C were measured and then mixed  by ball milling. Then,  the mixtures were  pressureless sintered at 2523 K for 3 h in an Ar atmosphere after metal  injection  molding process.  The oxidation tests were performed using a horizontal  tube furnace. Before the  tests, specimens were prepared using conventional polishing with a diamond abrasive, down to a 0.5 lm finish. They were then placed on an alumina boat and  inserted into the hot zone of the furnace. The oxidation tests were conducted at 1273  and 1473 K for 12 h under an air atmosphere.  312  Oxid Met (2016) 85:311-320  123  \\x0c', 'The densities of  the sintered specimens were measured using the Archimedes  method, and the theoretical densities of  the composites were calculated using the  rule of mixture. The microstructures of the cross-sections of the oxidized specimens  were  characterized  using scanning electron microscopy (SEM). To analyze  the  microstructures of  the vertical  sections after  the tests,  the specimens were cross sectioned and mounted in epoxy, carefully polished with a diamond abrasive down to a 0.5 lm finish, and cleaned in an ultrasonic bath. The thicknesses and element  profiles  of  the  resulting  reaction  layers were measured  and  analyzed  from the  polished cross-sections by energy dispersive spectrometer  (EDS).  Results and Discussion  Density and Microstructure  Table 1 shows the density of  the specimens obtained. The relative densities of  the  specimens before oxidation test  is in the range of 96-99 %.  Figure 1  shows  SEM images  of  the  polished  surfaces  of  the  ZrB2-SiC  composites. The darker phase is SiC, and it appears to be uniformly dispersed in  the lighter ZrB2 matrix. The microstructures of composites are regular, and little  pore was observed in the polished surfaces. Based on the high relative density and  the lack of any open pores, porosity should not have a significant effect on oxidation  behavior.  Oxidation Behavior at 1273 K  Thermodynamically, both ZrB2 and SiC should be oxidized when exposed to air.  However,  the oxidation rates of both species are negligible below about 1073 K  Table 1  Summary  of ZrB2-SiC specimens:  compositions,  designations,  bulk  densities,  theoretical  densities, and relative densities  Samples  Compositions  Bulk densities (g cm-3)  Theoretical densities (g cm-3)  Relative  densities (%)  ZrB2  (wt%)  SiC  (wt%)  B4C  (wt%)  ZrB2  99  0  1  5.79  6.00  96.5  ZrB2-  5%SiC  94  5  1  5.59  5.75  97.2  ZrB2-  10%SiC  89  10  1  5.37  5.52  97.3  ZrB2-  20%SiC  79  20  1  5.01  5.10  98.2  ZrB2-  30%SiC  69  30  1  4.65  4.75  97.9  Oxid Met (2016) 85:311-320  313  123  \\x0c', '314  Oxid Met (2016) 85:311-320  Fig. 1  Cross-sectional  images of a ZrB2-5%SiC, b ZrB2-10%SiC, and c ZrB2-30%SiC  [19]. Previous studies have reported that  the oxidation of ZrB2 by Reaction (1)  is  much faster  than that of SiC between 1073 and 1473 K [20, 21]. Assuming that  oxidation of ZrB2 proceeds stoichiometrically, temperature *723 K)  (melting  and  B2O3  the reaction should produce molten  solid ZrO2. Upon  cooling  to  room  temperature,  the B2O3 forms an amorphous solid while the ZrO2 is crystalline [18].  ZrB2 ðsÞ þ 5 2  O2 ðgÞ ¼ ZrO2 ðsÞ þ B2O3 ðlÞ  ð1Þ  Figure 2 shows  the SEM images of ZrB2-SiC specimens with different SiC  addition contents after exposure to air for 12 h at 1273 K. As shown in the figure,  the surface of oxide layer is fragile and easy to have crack and spallation, and it  is  easy to be exfoliated during cutting and polishing process.  In order  to observe the  oxide layer clearly,  the cross-section of ZrB2-SiC specimens without polishing are  analyzed by EDS. Figure 3 shows the EDS maps for ZrB2-10%SiC specimens after  oxidation at 1276 K in air for 12 h. The inserted table shows the mole ratio of the  elements C, O, Si and Zr of red mark point  in Fig. 3a. As the Ka of boron is 0.183,  Fig. 2  SEM images of cross section of ZrB2-SiC specimens after exposure to air a Without SiC addition, b ZrB2-5%SiC, c ZrB2-10%SiC, d ZrB2-20%SiC, e ZrB2-30%SiC  for 12 h at 1273 K.  123  \\x0c', 'Oxid Met (2016) 85:311-320  315  Fig. 3  EDS maps for ZrB2-10%SiC specimen after oxidation at 1273 K in air for 12 h. a SEM images of the cross section, b Zr map, c Si map, d C map, e O map. The inserted table is the analysis of the red  point  in (a)  which is lower than the detection limit of the equipment used,  the boron element  is  very difficult  to be detected. According to the mole ratio of O and Si,  it could be  found out  that  the glassy layer  is mainly composed by B2O3. Rezaie et al.  [19]  reported that  the surface structure consisted of 3 layers after the oxidation of ZrB2-  SiC at  1273 K for  30 min:  (1)  a  layer  of B2O3,  (2)  a ZrO2  layer  contained  unoxidized SiC,  (3)  unaffected ZrB2-SiC. Therefore,  a  continuous B2O3  layer  formed above the ZrO2 layer was found in the reported study. As shown in Fig. 3, a  Fig. 4  EDS maps for ZrB2-20%SiC specimen after oxidation at 1273 K in air for 12 h. a SEM images of the cross section, b Zr map, c Si map, d O map, e C map. The inserted tables are the analyses of the red  points in (a)  123  \\x0c', '316  Oxid Met (2016) 85:311-320  glassy layer of B2O3 is on the top of the surface. The maps of Zr, Si, C and O show  that a ZrO2 layer with SiC is under B2O3 layer, and the thickness of  this layer  is  larger  than that of B2O3 layer. Therefore,  the two-layer-structure was found in the  oxide layer of ZrB2-10%SiC specimen.  Figure 4  shows  EDS maps  of  ZrB2-20%SiC specimen  after  oxidation  at  1273 K and the  inserted tables  show the point  analysis of  the  red mark points.  Different with the ZrB2-10%SiC specimen,  the carbon element was not detected  on  the  top  oxide  layer,  which means  that  SiO2  exists  in  the  top  layer.  Additionally,  the  relatively higher  ratio of O indicate  the  composition of B2O3  should also exist  in the top layer. This means that a borosilicate glass is formed on  the top layer. The maps of Zr, Si, O and C show that ZrO2 with SiC is under  the  glassy layer, which is  same  as  that of  the ZrB2-10%SiC specimen. Therefore,  when SiC content  increased, some SiO2 is formed and the formation of SiO2 could  be accelerated by the presence of B2O3  [22].  Figure 5 shows  the thickness of  the oxidation layer of ZrB2-SiC composites.  With the increase of SiC addition content,  the oxide layer  thickness decreased.  It  was  reported that SiC additions do not  affect  the oxidation rate below 1373 K.  However,  in this work the increase of SiC addition shows an increase oxidation  resistance of ZrB2 at 1273 K. Seong et al. [23] reported that the oxidation kinetics of  ZrB2 might be  controlled by O2 diffusion and transport  through the ZrB2 grain  boundaries and ZrO2 grain boundaries, respectively. Although the oxidation of SiC  is much slower  than that of ZrB2  at 1273 K, when SiC content  is higher  than  10 wt% in our experimental conditions, SiO2 forms in the top glassy layer, which  could decrease the diffusion of oxygen.  Fig. 5  Thickness of oxide layer of ZrB2-SiC composites after exposure to air  for 12 h at 1273 K  123  \\x0c', 'Oxid Met (2016) 85:311-320  317  Oxidation Behavior at 1473 K  When ZrB2-SiC was heated at 1473 K, the composition and structure of the surface  layers changed. As the temperature approaches 1473 K, the vapor pressure of B2O3  increases,  leading to its  rapid evaporation.  In addition, SiC starts  to oxidize  to  produce SiO2  and CO. According  to Shugart  et  al.  [24]  study, B2O3 does not  completely evaporate even at 1773 K,  thus little B2O3 should be still  remained in  the oxide  layer. Figure 6 shows  the SEM images of ZrB2-SiC specimens with  different SiC addition contents after exposure to air for 12 h at 1473 K. An apparent  oxide layer  is formed for ZrB2-SiC compositions. Figure 6d shows that  the oxide  layer is not as compact as the unaffected ZrB2-SiC and some pores are visible in the  oxide layer. When SiC addition content is 5 wt%, a two-layer-structure is formed as  follows: (1) a layer of ZrO2 and SiO2 with unreacted SiC; (2) unaffected ZrB2-SiC.  As a little boron is detected in oxide layer,  it  is considered that a small amount of  B2O3  is  still  remain in the oxide layer. However, when SiC addition content  is  higher  than 10 wt%, a continuous molten layer could be observed. As the melting  point of SiO2 is higher than 1700 K, a borosilicate glass layer is formed on the top  of  the oxide layer.  Fig. 6  SEM images of cross section of ZrB2-SiC specimens after exposure to air a ZrB2-5%SiC, b ZrB2-10%SiC, c ZrB2-20%SiC, d ZrB2-30%SiC  for 12 h at 1473 K.  123  \\x0c', '318  Oxid Met (2016) 85:311-320  Fig. 7  SEM images of cross section of ZrB2-SiC specimens after exposure to air a ZrB2-30%SiC, b high magnification of square area of (a)  for 12 h at 1473 K.  Fig. 8  Thickness of oxide layer of ZrB2-SiC composites after exposure to air  for 12 h at 1473 K  According to the reported TEM study of the oxide layer on ZrB2-SiC [23], ZrB2  was oxidized and transformed to ZrO2 phase firstly and then SiC was oxidized at the  interface  between  unreacted  layer  and  oxidized  layer  during  the  oxidation.  Figure 7b shows  that SiC started to oxidize  and transform into SiO2  from the  surface of SiC grain. After that, the SiO2 was dispersed in grain boundaries in whole  oxide  layer  on  composite  due  to  high  viscosity  and  volumetric  increase. The  unreacted SiC existed in amorphous SiO2.  Figure 8 shows the thickness of the oxidation layer of ZrB2-SiC composites. The  addition of SiC greatly decreases  the oxide layer  thickness because of  the SiO2  formation. When SiC content  is 5 wt%,  the oxide layer thickness is almost half of  that without SiC addition. When the SiC content increased from 5 % to 30 wt%, the  oxide  layer  thickness decreased gently. This  result proved that  increase of SiC  addition  could  effectively  improve  the  oxidation  resistance  of ZrB2.  In  all  the  specimens, ZrB2 containing 30 wt% SiC exhibited the highest oxidation resistance  at 1473 K in air.  123  \\x0c', 'Oxid Met (2016) 85:311-320  319  Conclusions  The oxidation behavior of ZrB2-SiC composites with different SiC contents was  investigated at 1273 and 1473 K in air for 12 h. The following conclusions could be  obtained:  (1)  When ZrB2-SiC composite is oxidized at 1273 K in air, a two-oxide layer structure  forms:  a  continuous  glassy  layer  and  a ZrO2  layer  contained  unoxidized SiC. When SiC content  is 5 and 10 wt%,  the glassy layer  mainly  composed  by B2O3. When  SiC content  is  20  and  30 wt%,  is  a  borosilicate glass could be formed on the top layer, which could improve the  oxidation resistance of ZrB2.  (2)  When ZrB2-SiC composite is oxidized at 1473 K in air,  the oxide layer  is  composed of ZrO2 and SiO2 with unreacted SiC. When SiC addition content  is higher than 10 wt%, a continuous borosilicate glass layer could be formed  on the top of  the oxide layer.  (3)  With  the  increase  of SiC content  in ZrB2,  the  thickness  of  oxide  layer  decreased at both 1273 and 1473 K. The addition of SiC shows effective  protection of ZrB2 at 1473 K.  (4)  In  all  the  specimens, ZrB2  containing  30 wt% SiC exhibits  the  highest  oxidation resistance at 1473 K in air.  References  1. M. M. Opeka, I. G. Talmy and J. A. Zaykoski, Journal of Materials Science 39, 5887-5904 (2004).  2. F. Monteverde, A. Bellosi and S. Guicciardi, Journal of  the European Ceramic Society 22, 278-288  (2002).  3. R. A. Cutler, in Engineering Properties of Borides, ed. S. J. Schneider, (ASM International, Materials  Park, 1991).  4. W. C. Tripp and H. C. Graham, Journal of  the Electrochemical Society 118, 1195-1209 (1971).  5. W. G. Fahrenholtz, Journal of  the American Ceramic Society 88, 3509-3512 (2005).  6. S. R. Levine, E.  J. Opila, M. C. Halbig,  J. D. Kiser, M. Singh and J. A. Salem, Journal of  the  European Ceramic Society 22, 2757-2767 (2002).  7. A. K. Kuriakose and J. L. Margrave, Journal of  the Electrochemical Society 111, 827-831 (1964).  8. E. J. Opila and M. C. Halbig, Ceramic Engineering & Science Proceedings 22, 221-228 (2001).  9. A. L. Chamberlain, W. G. Fahrenholtz and G. E. Hilmas, Refractories Applications Transactions 1,  1-8 (2005).  10. F. Monteverde and A. Bellosi, Journal of  the Electrochemical Society 150, B552-B569 (2003).  11. D. Sciti, M. Brach and A. Bellosi, Journal of Materials Research 20, 922-930 (2005).  12. W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy and J. A. Zaykoski, Journal of  the American Ceramic  Society 90, 1347-1364 (2007).  13. S. S. Hwang, A. L. Vasiliev and N. P. Padture, Materials Science and Engineering: A 464, 216-224  (2007).  14. P. Hu, W. Guolin and Z. Wang, Corrosion Science 51, 2724-2732 (2009).  15. W. G. Fahrenholtz, Journal of  the American Ceramic Society 90, 143-148 (2007).  16. A. Rezaie, W. G. Fahrenholtz and G. E. Hilmas, Journal of  the American Ceramic Society 89,  3240-3245 (2006).  123  \\x0c', '320  Oxid Met (2016) 85:311-320  17. P. A. Williams, R. Sakidja, J. H. Perepezko and P. Ritt, Journal of the European Ceramic Society 32,  3875-3883 (2012).  18. W. Han, P. Hu, X. Zhang,  J. Hang and S. Meng, Journal of  the American Ceramic Society 91,  3328-3334 (2008).  19. A. Rezaie, W. G. Fahrenholtz and G. E. Hilmas, Journal of  the European Ceramic Society 27,  2495-2501 (2007).  20. W. C. Tripp, H. H. Davis and H. C. Graham, American Ceramic Society Bulletin 52, 612-613 (1973).  21. W. C. Tripp and H. C. Graham, Journal of  the Electrochemical Society 118, 1195-1199 (1968).  22. C. E. Ramberg, G. Cruciani, K. E. Spear, R. E. Tressler and C. F. Ramberg, Journal of the American  Ceramic Society 79, 2897-2911 (1996).  23. Y. H. Seong, S.  J. Lee and D. K. Kim, Journal of  the American Ceramic Society 96, 1570-1576  (2013).  24. K. Shugart, S. Liu, F. Craven and E. Opila, Journal of  the American Ceramic Society 98, 287-295  (2015).  123  \\x0c']"
},{
  "_id": 51,
  "PDF": "Effect of SiC content on electrical, thermal and ablative properties of pressureless sintered ZrB 2 -based ultrahigh temperature ceramic composites.pdf",
  "Text": "['Journal of the European Ceramic Society 37 (2017) 559-572  Contents lists available at www.sciencedirect.com  Journal of the European Ceramic Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Feature article  Effect of SiC content on electrical, thermal and ablative properties of pressureless sintered ZrB2-based ultrahigh temperature ceramic composites  Manab Mallik a , Ansu. J. Kailath b , K.K. Ray c , R. Mitra c,∗  a Department of Metallurgical and Materials Engineering, National Institute of Technology, Durgapur 713209, West Bengal, India b National Metallurgical Laboratory, Jamshedpur 831007, Jharkhand, India c Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, West Bengal, India  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 31 May 2016 Received in revised form 2 September 2016 Accepted 23 September 2016 Available online 1 October 2016  Keywords:  Borides Composites Electrical conductivity Thermal properties Ablation resistance  1.  Introduction  The effects of SiC content (10-40 vol.%) on electrical, thermal and ablation properties of pressureless sintered ZrB2 -SiC composites showing interfacial segregation of W-rich phases have been studied. The electrical resistivity was measured by four-probe method, whereas thermal diffusivity and coefficient of thermal expansion (CTE) were determined using laser-flash method and thermo-mechanical analyzer, respectively. Whereas thermal conductivities calculated from experimentally obtained thermal diffusivity values are found to be the highest for the ZrB2 -20 SiC composite, both electrical conductivity and CTE decrease with increasing SiC content. The specimens were subjected to thermal shock by soaking at 800-1200   C, followed by water-quenching. Further, some specimens were exposed to oxyacetylene flame (2200   C) for 10 min. The damage was estimated from changes in mass, Young’s modulus, and hardness. The highest thermal shock and ablation resistance have been observed for the ZrB2 -20 SiC composite, as thermal properties and formation of protective oxide scale play key role. © 2016 Elsevier Ltd. All rights reserved.  The hypersonic flight vehicles and reusable atmospheric reentry vehicles are designed to possess sharp leading edges (radius as low as 1 mm) and smooth surfaces in order to improve their aerodynamic performances [1]. But, the sharp surfaces encounter extraordinary thermal and oxidative loads due to aerothermal heating during re-entry, when the surface temperature is found to be proportional to the inverse square root of the body radius. Therefore, it can be said that smaller the radius of a sharp leading edge or nose cone, higher is the surface temperature, which causes severe ablation or blunting of the available thermal protection system materials. Therefore, the materials for use in the thermal protection system (TPS) should resist thermal shock and ablation for reliable operation over a multi-mission life-cycle [2]. The refractory metal diborides have high melting points (>3000  C) and ability to form protective oxide scales, which provide capability of operating in oxidizing environments for short  ∗ Corresponding author. E-mail addresses: rahul@metal.iitkgp.ernet.in, rahulmitra1966@gmail.com (R. Mitra).  http://dx.doi.org/10.1016/j.jeurceramsoc.2016.09.024 0955-2219/© 2016 Elsevier Ltd. All rights reserved.  durations at temperatures of 2000  C or above [3-5]. Among all the UHTCs, ZrB2 has the lowest theoretical density (6.09 g/cm3 ), which combined with its high melting temperature, excellent resistance to thermal shock and oxidation as well as high electrical and thermal conductivities is responsible for triggering extensive research on this material. [6-10]. Such impressive properties of ZrB2 make it attractive for use in hyersonic re-entry vehicles requiring brief exposures at temperatures ≥2000  C [7,8]. The results of a few earlier studies have shown that addition of SiC as reinforcement to ZrB2 based composites improves their sinterability, as well as leads to considerable increase in their strength, toughness and oxidation resistance [11-18]. Moreover, it has been reported by Zhang et al. that the ZrB2 -20 vol.% SiC composite exhibits superior thermal-oxidative and configurational stability in the simulated re-entry environment compared to that of the C/SiC composite [19]. Zhou et al. have investigated the ablation behavior of ZrB2 -SiC-ZrO2 ceramics, and have reported that excessive amount of ZrO2 provides path for inward transport of oxygen flow [20]. Further, a study by Dehdashti et al. [21,22] on the effects of transition metals such as W, Mo and Nb incorporation on oxidation behavior of ZrB2 , show W to be least effective at improving the oxidation resistance of ZrB2 at 1600  C.  \\x0c', '560  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  The results of previous investigations have shown that WC addition has beneficial effect on densification and mechanical properties of the ZrB2 -SiC based composites [23-25]. Effect of WC addition on enhancement of densification of ZrB2 -SiC composite has been reported by Zou et al. It has been suggested that liquid phase formation by melting of (W, Zr)ss B and (W, Zr)ss Si2 leads to the improvement of densification [24]. Furthermore, it has been shown that addition of 5 vol.% WC to the ZrB2 -20 vol.% SiC composite improves the high temperature flexural strength, when the tests are carried out in inert environment [25]. Significant improvement in the strength at high temperature is caused by removal of the oxide impurities from grain boundaries. Generally, high thermal conductivity improves the thermal shock resistance by reducing thermal gradients within the material, and also allows the heat energy to be conducted and radiated away from the surface. Monolithic ZrB2 has thermal conductivity values in the range of 56-83.8 Wm−1 K−1 at room temperature [26,27], whereas the thermal conductivity values of SiC range from 100 to 350 Wm−1 K−1 [28]. It has been reported that addition of 20 and 30 conductivity by ≈18% and ≈11%, respectively [26,27]. On the other vol.% SiC as reinforcement to ZrB2 leads to increase in its thermal hand, studies on the effects of transition-metal (hafnium, niobium, tungsten, titanium, and yttrium) additions to ZrB2 as reported by McClane et al. [29] have shown significant reduction in thermal conductivity. The thermal strains imposed in the ceramic components subjected to high temperature exposure depend on their coefficient of thermal expansion (CTE), and therefore their thermal shock (5.9 × 10−6 K−1 resistance is also affected. The CTE of ZrB2 [10]) in the temperature range of 27-1027  C is higher than that of SiC (4.3 × 10−6 K−1 [30]). Therefore, addition of SiC as reinforcement is expected to lower the CTE of ZrB2 -SiC composites, and thereby improve their dimensional stability. It should also be noted that the matrix grain size, as well as matrix-reinforcement interfacial area and bond strength have a very strong influence on the CTE of the composites [31-33]. In an earlier study, it has been shown that with increase in SiC content of the ZrB2 -SiC composites pressureless-sintered with B4C and C as additives, densification, hardness and Young’s modulus are increased, whereas matrix grain size is reduced [16]. It has been shown that contamination caused by wear of WC-Co balls and vials during mixing leads to the formation of W-rich phase at both ZrB2 SiC interfaces and matrix grain boundaries. Further, the amount of W-contamination has been found to scale with the SiC content, because of its abrasive action during ball-milling. The purpose of this study is to investigate the effect of SiC content and interfacial W-rich phase, if any on coefficient of thermal expansion, electrical and thermal conductivity as well as thermal shock and ablation resistance of the pressureless sintered ZrB2 -SiC composites.  2. Experimental procedure  Commercially available powders of ZrB2 , SiC and B4C, supplied by H. C. Starck Ltd. Germany, were used as raw materials in this study. The powder samples having composition of ZrB2 -10 vol.% SiC-5.6 vol.% B4C and 4.8 vol.% C (ZSBC-10), ZrB2 -20 vol.% SiC-5.6 vol.% B4C and 4.8 vol.% C (ZSBC-20), ZrB2 -30 vol.% SiC-5.6 vol.% B4C and 4.8 vol.% C (ZSBC-30), and ZrB2 -40 vol.% SiC-5.6 vol.% B4C and 4.8 vol.% C (ZSBC-40) were mixed in aplanetary mono-mill in ethanol at 250 rpm for 6 h using WC-Co media inside WC-Co vials. Here, phenolic resin was added as a binder for preparation of green bodies, and also as the source of carbon, to be left as residue during heating for sintering. After milling, the powders were dried and crushed using an agate mortar and pestle. The composite powders were compacted into pellets by uniaxial pressing at 100 MPa using  a 25 mm diameter steel die. Subsequently, the pellets were sintered at 2000  C for 2 h in a graphite resistance-heating ASTRO furnace. The furnace was heated from the ambient temperature at the rate of 10  C min−1 , and before reaching the sintering temperature, the green compacts were held for 1 h at 850  C to convert the resin into carbon, and subsequently for 1/2 h at intermediate temperatures of both 1250  C and 1600  C to ensure the reduction of surface oxide impurities by B4C and C [16], which can be predicted on the basis of thermodynamic data [34,35]. Archimedes(cid:6) principle was utilized to assess the bulk density of each pressureless sintered sample. After proper sectioning of the sintered samples by a slow speed diamond saw, those were prepared for microstructural observation by metallographic polishing. X-Ray diffraction (XRD) analysis was utilized to identify the different phases in the microstructures by using Co K␣ radiation. The microstructures of the sintered composites were examined by field emission scanning electron microscopy (FESEM, SUPRA 40, CARL ZEISS SMT Ltd.), and the chemical compositions were simultaneously determined using energy dispersive X-Ray (EDX) analyzer. The specimens for TEM analysis were prepared using mechanical thinning followed by ion-beam milling. Subsequently, the microstructures of the specimens were studied on a JEM 2100UHR (JEOL, Japan) transmission electron microscope operated at an acceleration voltage of 200 kV. Both bright field and dark field imaging modes were used to investigate the microstructure of the samples while the structures of the constituent phases were evaluated using selected area electron diffraction (SAED). Image analyses software has been used to estimate average grain size that was obtained by measuring more than 100 grain dimensions using both SEM and TEM images. The sintered composites were sliced using wire electrodischarge machining (EDM) for electrical resistivity and thermal diffusivity measurements. The van-der-Pauw four-probe method equipped with a high precision resistivity unit (Keithley 2182A nanovoltmeter and 6220 current source) has been used to measure the electrical resistivities of the investigated composites at room temperature. Specimens having diameter of 20 mm and thickness in the range of 2-4 mm were used to measure the electrical resistivity (Rs ) using the following relation:  RS = (cid:2)tR  ln 2  (1)  where t is the thickness and R is the resistance of the specimen. The laser flash method was utilized to measure the thermal diffusivity of the of ZrB2 -SiC composites according to the ASTM E-1461 standard [36]. Disc shaped specimens (having dimensions − 12.7 mm diameter and 1.0 mm thickness) were exposed toa radiant energy pulse by focusing a high intensity laser beam for a short duration. The increase in the rear-face temperature of the specimen due to the incident energy absorbed on the front surface was monitored to record the time required for the temperature to reach 50% of the maximum value (t1/2 ). This time was recorded continuously with a purpose to calculate the thermal diffusivity (a) using the relation [36]:  ˛ = 0.13879 L2  t1/2  (2)  where L is the specimen thickness and t1/2 is the half time. The specific heat (c) of each investigated composite was evaluated using the rule of mixtures from the specific heat data available in the literature [36] for ZrB2 , SiC, B4C and W by considering their weight fractions. Thereafter, the thermal conductivity (\\u242d) was calculated using the relation [36]:  \\u242d = ca(cid:3)  (3)  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  561  where \\u2433 is the density of the material measured at room temperature (25  C). A thermo-mechanical analyzer (TMA, Model: Diamond TMA, Perkin Elmer, USA) was utilized to measure the coefficients of thermal expansion (CTE) of the ZrB2 -SiC composites in the temperature range of 200-1000  C. For this purpose, specimens with dimensions of 10 mm (length) × 5 mm (width) × 4 mm (thickness) were heated at the rate of 5  C per minute in argon atmosphere, and the changes in length with increase in temperature were continuously recorded. The value of CTE (␣) was calculated using the relationship:  ˛ =  (cid:4)L  L0(cid:4)T  (4)  where Lo is the length of the samples at ambient temperature, To 20  C, (cid:4)L and (cid:4)T are the changes in the length and the temperature, respectively with respect to their initial values. Thermal shock tests were conducted on small specimens having the dimensions of 10 mm × 6 mm × 4 mm. The samples were first ground and polished metallographically. Subsequently, the polished surfaces were indented using Vickers indenter operated at 20 kgf load, and the lengths of median cracks formed at the indentation corners were measured by FESEM. These samples were soaked at 800  C, 1000  C, or 1200  C for 10 min in air by hanging inside a vertical chamber furnace using nichrome wires. Subsequently, the samples were taken out of furnace and quenched in water. The damage due to thermal shock was assessed quantitatively by measuring the relative changes in hardness and indentation crack lengths. The relative ablative behaviors of the investigated composites were studied by subjecting the specimens with dimensions of 20 mm diameter and 4 mm thickness, embedded in a graphite block to an oxyacetylene gun with inner diameter of 2 mm. The surface temperatures of the samples were measured at intervals of every 15 s. An optical pyrometer was used to measure the temperature at the front face, while a thermocouple was used to measure the temperature of the back face. A neutral flame was used to achieve the target gas temperature of 2000-2200  C. Each specimen was subsequently exposed for 10 min duration. For a comparative assessment of the resistance to high temperature ablation, the composite samples with 20 mm diameter were held inside a groove made in a graphite holder, and were then exposed to a neutral oxy-acetylene flame for 10 min. The front-face temperature was continuously monitored using an optical pyrometer, whereas the back-face temperature was measured using a W-Re thermocouple. The experimental set-up for the ablation experiment was designed on the basis of the procedure used in an earlier study [37]. The specimens were photographed and weighed before and after the tests. The damage in the samples exposed to oxy-acetylene flame was quantified in terms of the changes in mass, and Young’s modulus. Phase identification by XRD analyses as well as examination of chemical compositions through EDX analyses on SEM were used to identify the constituents of the oxide scale formed by elevated temperature exposure.  3. Results and discussion  3.1. Microstructure  The densities of the ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites have been found to be 5.62, 5.42, 5.16 and 4.92 g/cm3 , respectively. By comparing these experimentally measured density values with theoretical densities calculated using the rule of mixtures, the relative densities of ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 are found as 98%, 99.5%, 99.6% and 99.6%, respectively. The typical SEM (BSE) images of the ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites are shown in Fig. 1(a) through (d), respectively.  The constituent phases, ZrB2 , SiC and B4C can be conveniently distinguished from one another in the SEM (BSE) images on the basis of differences in their atomic number contrast and the results of EDX analysis. No evidence of excess carbon has been found in the microstructures of the investigated pressureless-sintered composites, indicating that its amount is negligible. This can be attributed to near-complete consumption of C in reduction of the oxide impurities. Based on the analysis of the oxygen contents of the raw materials (as mentioned in Ref. 16), it is found that the oxide impurities would be more or less fully reduced by the added carbon. The grain boundaries of ZrB2 and ZrB2 -SiC interfaces show clear evidence of segregation of W, Fe, and Co, which appear to have promoted densification by formation of a liquid phase during sintering. The amounts of W present as impurity in the sintered composite pellets, and measured by bulk EDX analyses have been found as 2.6 ± 0.6, 3.5 ± 0.9, 4.2 ± 0.8, and 7.7 ± 0.4 wt.% in ZSBC-10, ZSBC-20, ZSBC-30, and ZSBC-40 composites, respectively [16]. Furthermore, the concentrations of both Co and Fe have been found to vary in the range of 0.1-0.3 wt.%. Whereas the impurities like W and Co are believed to be contributed by erosion of WC-Co vial walls and milling media, Fe was probably added through wear and tear of the die walls and ram made of die-steel during cold compaction of the powder mixtures to prepare pellets for sintering. Furthermore, the average grain size is decreased from 23.1 \\u242em to 11.7 \\u242em with increase in SiC content from 10 to 40% [16]. This suggests that the amount of grain-boundary or interfacial segregation of W-rich phase increases with increase in volume fraction of SiC, which has been also effective in inhibiting grain-growth. Typical bright field TEM images depicting a typical ZrB2 matrix grain boundary-SiC triple point interface along with SAED pattern is shown in Fig. 2(a) and (b), respectively, and this location exhibits the enrichment of W and Co as confirmed by EDX analysis (Fig. 2(c)). Analysis of the SAED pattern confirms the formation of Co0.9W0.1 phase with hexagonal crystal structure.  3.2.  Electrical resistivity and conductivity  The isotropic nature of electrical properties in the investigated composites was confirmed from the resistances measured on the ZrB2 -SiC composite samples between the pairs of mutually orthogonal contacts, which were found to be more or less similar. The electrical resistivity and conductivity values obtained by using the Van-der-Pauw four-probe method [38] (shown in Table 1) indicate that the electrical resistivities of the investigated composites increase in the order of ZSBC-10 < ZSBC-20 < ZSBC-30 < ZSBC-40. The results obtained for the hot-pressed ZrB2 -20 vol.% SiC composites have been included in Table 1 for the purpose of comparison [39]. The combined influence of increasing volume fraction of SiC and particle-matrix interfacial area contributes to the increase in the electrical resistivity. In an earlier study, the electrical resistivity has been found to increase by ≈2.5 times on addition of 20 vol.% SiC particles to the ZrB2 matrix [40]. Interestingly, the reported values of the electrical resistivity of ZrB2 have been found to vary in the range of 6.7-22 \\u242e(cid:5) cm [26,41,42], which can be attributed to differences in matrix grain sizes and purity of raw materials. The Brick Layer Model (BLM) has been employed to estimate the average values of internal interfacial electrical resistance for the investigated ZrB2 -SiC composites [43]: ␥ = [(1/␥int ) + (R I /␦)]  (5)  −1  where ␥ is the electrical conductivity of a typical single phase material (the matrix phase, ZrB2 in this study), ␥int is the intrinsic electrical conductivity of the matrix phase, RI is the average interfacial electrical resistance, and ␦ is the average matrix grain size. Further, the value of ␥ used in Eq. (5) has been calculated using the effective medium approximation (EMA), by considering  \\x0c', '562  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  Fig. 1. SEM (BSE) images depicting the microstructures of (a) ZSBC-10, (b) ZSBC-20, (c) ZSBC-30 and (d) ZSBC-40 composites.  Fig. 2. (a) Bright field TEM image depicting matrix grain boundary-particle triple point interface, along with (b) SAED pattern, and (c) EDX spectrum from the Co0.9W0.1 solid solution phase.  Table 1 Estimated electrical resistivities and conductivitis of the investigated composites. The results of hot-pressed ZrB2 -SiC (ZS) composites are shown for comparison [39].  Materials  Specimen Thickness (mm)  Voltage (mV)  Current (mA)  Resistance (\\u242e(cid:5))  Resistivity (10−8 (cid:5)-m)  ZS ZSBC-10 ZSBC-20 ZSBC-30 ZSBC-40  1.7 2 1.6 2 2  0.00133 0.00234 0.00399 0.00598 0.00980  100 100 100 100 100  13.3 23.4 39.9 59.8 98.0  10.25 21.00 28.64 53.66 87.93  Conductivity (106 S/m) Measured  9.76 4.8 3.5 1.86 1.14  the SiC particles as isolated inclusions surrounded by ZrB2 medium with an effective electrical conductivity, ␥eff . The value of ␥eff , i.e. the experimentally obtained electrical conductivity for each of the  investigated composites is related to the intrinsic electrical conductivities of ZrB2 and SiC through the following relation [44]:  N(cid:2)  i=1  (cid:6)eff − (cid:6)i  2(cid:6)eff + (cid:6)i  fi = 0  (6)  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  563  where ␥i and fi are the intrinsic electrical conductivity and the volume fraction of the ith phase, respectively. For these calculations, the intrinsic electrical resistivities of ZrB2 and SiC have been taken as 4.9 \\u242e(cid:5)-cm [45] and 5 × 103 \\u242e(cid:5)-cm [46], respectively. The values of the electrical conductivities of the matrix ␥i (=␥m ) have been evaluated using Eq. (6) for each of the investigated composites. Owing to the polycrystallinity of the matrix with the electrical resistances being offered by the grain boundaries, its intrinsic conductivity cannot be considered to be same as that of a single crystal. This electrical conductivity of ZrB2 (␥m ) as obtained from Eq. (6) has been considered as the value of ␥ in Eq. (5), similar to that followed by Zhang et al. [40]. The values of RI calculated for ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites have been found as 3.02 × 10−12 m2 /S, 3.2 × 10−12 m2 /S, 3.5 × 10−12 m2 /S, and 3.7 × 10−12 m2 /S, respectively. An increasing volume fraction of SiC is observed to affect the values of RI marginally, as this reinforcement phase possesses lower electrical conductivity than that of the ZrB2 matrix. It is intuitive that the increase in the net interfacial area with SiC volume fraction due to the increasing amount of W-rich phase at ZrB2 matrix grain boundaries and particle-matrix interfaces would also contribute to the increase in interfacial electrical resistance of the investigated composites. Interestingly, the results in Table 1 show that the electrical resistivity of the ZSBC-20 composite is found to be 2.8 times higher than that of the hot-pressed ZrB2 -20 vol.% SiC (ZS) composite [39]. On the other hand, the value of RI (1.5 × 10−13 m2 /S) reported for ZS is less by an order of magnitude compared to that of the ZSBC-20, in spite of same volume fraction of SiC particles, comparable relative density and finer matrix grain size of the former composite. This observation can be explained only on the basis of the interfacial segregation of W-rich phase, which probably increases the contribution of the interfacial electrical resistance in the pressureless sintered ZSBC-20 composite in comparison to that in the ZS, which has shown sharp ZrB2 -SiC interfaces without significant impurity segregation [17,39].  3.3.  Thermal diffusivity, specific heat and thermal conductivity  Plots depicting the temperature dependence of thermal diffusivity, specific heat and thermal conductivity of the pressureless sintered ZrB2 -SiC composites are shown in Fig. 3(a)-(c), respectively. The thermal conductivity has been calculated from the experimentally obtained thermal diffusivity data using Eq. (3). For this purpose, the density measured at the room temperature (25  C) using Archimedes(cid:6) principle has been used. The specific heat values of the developed composites at different temperatures (shown in Fig. 3(b)) have been calculated on the basis of the rule of mixtures by considering weight fraction of matrix and reinforcement phases, and the literature-based data for individual phases [35]. It is found that the ZSBC-20 composite exhibits the highest thermal the maximum thermal diffusivity value of ≈ 0.26 cm2 /s at diffusivity among all the presently investigated composites with room temperature. A similar trend is observed at all temperatures up to 1200  C (Fig. 3(a)). The thermal diffusivity of the ZrB2 -SiC composites at a given temperature is expected to be affected by both electronic and phonon-based thermal transport mechanisms [26,27,39,40]. Zhang et al. [40] have calculated the ratio of electronic component (\\u242de ) to total thermal conductivity (\\u242d) at room temperature (25  C) using the relation:  \\u242de /\\u242d = 298L0 /\\u2433.\\u242d(7)  where \\u2433 is the electrical resistivity and L0 is the Lorenz number, whose value has been taken as = 2.45 × 10−8 W((cid:5) K2 )−1 by Zhang  Fig. 3. Plots depicting the variation of: (a) thermal diffusivity, (b) specific heat and (c) thermal conductivity with temperature for pressureless sintered ZrB2 -SiC composites. The thermal conductivity values of the hot-pressed ZrB2 -SiC composite (ZS) is shown in (c) for the purpose of comparison [39].  et al. [40]. The Eq. (7) has been used to estimate the ratio of \\u242de /\\u242d for the investigated ZrB2 -SiC composites, and the calculated values are summarized in Table 2. The results mentioned in this table show that the electronic contribution to thermal conductivity decreases with increasing SiC content or decreasing ZrB2 content of the composites, which is in agreement with the observations of Patel et al. [47]. Hence, based on the aforementioned results and analyses, it is plausible that the thermal transport is facilitated by electrons in the ZrB2 phase and drawn from the study of Zhang et al. showing value of \\u242de /\\u242d ≈ 1.0 by phonons in the covalently bonded SiC. Same inference can be for pure ZrB2 [40].  \\x0c', '564  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  Table 2 The fractions of contributions to total thermal conductivity (\\u242d) at room temperature (25   C) by phonons (\\u242dph ) and electrons (\\u242de ). The values obtained for hot-pressed ZrB2 -SiC (ZS) composites are shown for comparison [39].  Composites  \\u242dph (W/mK)  \\u242de (W/mK)  ZS ZSBC-10 ZSBC-20 ZSBC-30 ZSBC-40  18.30 19.34 38.58 48.40 42.40  71.23 34.77 25.72 13.70 8.10  \\u242de /\\u242d  0.80 0.64 0.40 0.22 0.16  A modified version of Eq. (5) has been used to calculate the average interfacial thermal resistances (R\\u242d ) for the ZSBC composites, with \\u2434int and RI being substituted by \\u242dint and R\\u242d , respectively, where \\u242dint is the intrinsic thermal conductivity of the matrix phase, and R\\u242d is the interfacial thermal resistance. By considering \\u242dint ≈ 140 W/m K for ZrB2 [48] and average matrix grain sizes of ZSBC-10, ZSBC-20 and ZSBC-30 as 23.1 \\u242em, 20.6 \\u242em and 14 \\u242em, respectively, the R\\u242d values of ZSBC-10, ZSBC-20 and ZSBC-30 composites are calculated as 4.0 × 10−7 m2KW−1 , 4.3 × 10−7 m2KW−1 and 6.1 × 10−7 m2KW−1 , respectively. These values lie well-within a range (10−7 -10−9 m2 KW−1 ), which is considered as typical for polycrystalline ceramics, as reported in the previous studies [40]. The value of R\\u242d observed for the investigated ZSBC-20 composite is 25.3 times greater than that reported for the hot-pressed ZrB2 -20 vol.% SiC composite (R\\u242d≈ 1.7 × 10−8 m2KW−1 ) by Mallik et al. [39]. This difference in R\\u242d in spite of finer grain size of the hot-pressed composite may be attributed to the presence of W-rich interfacial phase in the pressureless sintered composite, in contrast to the sharp particle-matrix interfaces in the former material. Furthermore, the thermal conductivity values found in the present study at both ambient and elevated temperatures are 24-72% less than those reported earlier [26,31,39], which is consistent with the difference between the values of R\\u242d , and is also due to the presence of W-rich interfacial phase in the pressureless sintered composites. As shown in Fig. 3(c), the difference between the experimentally determined thermal conductivities of the hot-pressed ZrB2 -SiC [39] and ZSBC-20 composites is 72.5% at 100  C, and is 28.8% at 1200  C. The decrease in the difference between the thermal conductivity values of hot-pressed and pressureless sintered composites with increasing temperature is attributed to increased contribution of phonon-scattering. It is interesting to note that in spite of higher R\\u242d , the ZSBC-20 composite shows higher thermal diffusivity and conductivity compared to the ZSBC-10, but the difference decreases with increasing temperature (Fig. 3). The results in Fig. 3 also show that the thermal conductivity of the ZSBC-10 composite is less than that of ZSBC-30 at temperatures <900  C, in spite of higher R\\u242d of the latter composite. These observations could be attributed to lower relative density of the ZSBC-10 (≈98%) compared to that of ZSBC-20 (≈99.5%) or ZSBC-30 (≈99.6%). At the ambient temperature, the ratio of thermal conductivity (value for single crystal) of SiC to that of ZrB2 is ≈1.36, whereas it is ≈0.42 at 1000  C [27,49]. Therefore, another reason for higher thermal conductivity of the ZSBC-20 and ZSBC-30 composite at lower temperatures can be their higher SiC volume fraction compared to that of the ZSBC-10. Interestingly, the thermal conductivity of the ZSBC-10 composite is less than that of the ZSBC-20 even at higher temperatures, at which SiC has much lower thermal conductivity compared to that of ZrB2 , suggesting that the effect of relative density or porosity content is probably more significant. Further, the amount of SiC in the ZSBC-20 composite is not enough to form a near-continuous network as in the composites with higher SiC volume fraction [16], and therefore the thermal transport is primarily through the ZrB2 matrix. The observed decrease in difference between the thermal conductivities of ZSBC-10 and ZSBC-20 composites with increasing temperature could be attributed to higher  Fig. 4. Thermal expansion curves for ZBC, ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 for the temperature range of 200   C to 1350   C.  R\\u242d of the latter material. The thermal conductivity of ZSBC-30 composite is less than that of ZSBC-10 at temperatures >900  C, whereas it is less than that of ZSBC-20 at all temperatures, which can be ascribed to the role of interfacial thermal resistance as well lower thermal conductivity of SiC than that of ZrB2 at higher temperatures. It is interesting to note that thermal conductivity decreases with increase in SiC content beyond 20 vol.% even at lower temperatures, in spite of higher thermal conductivity of SiC than that of ZrB2 . This observation is attributed to increase in interfacial resistance caused by larger matrix grain boundary area caused by decrease in grain size as well as increase in W-rich interfacial phase. Furthermore, it may be inferred that heat conduction in the composites with higher SiC volume fraction is affected by enhanced scattering of phonons in SiC due to the presence of ZrB2 grain-boundaries and ZrB2 -SiC interfaces. Based on the aforementioned observations, it is probably appropriate to infer that the role of interfacial thermal resistance in the investigated pressureless-sintered composites becomes significant beyond a critical SiC volume fraction. In a few earlier studies, the thermal conductivity of hotpressed ZrB2 -30 vol.% SiC composite has been found to be marginally higher than that of ZrB2 -20 vol.% SiC composite at the ambient temperature, but the trend is reversed at temperatures >200  C [26,47]. Although the thermal diffusivity of SiC is higher than that of ZrB2 at room temperature, yet it decreases faster with increasing temperature. For example, the reported values of thermal diffusivities of ZrB2 and SiC at room temperature are 0.225 cm2 /s and 0.50 cm2 /s, respectively, whereas these are 0.145 cm2 /s and 0.092 cm2 /s, respectively at 1000  C [26,49]. As the thermal diffusivity of SiC depends primarily on phononphonon, phonon-defect, and phonon-grain boundary interactions, it decreases with increasing temperature due to the increase in the amount of phonon-scattering [26]. It may be therefore inferred that the composites with higher SiC content tend to show sharper decrease in thermal conductivity with increasing temperature.  3.4.  Coefficient of thermal expansion (CTE) of ZSBC composites  The variation of linear thermal expansion per unit length (\\x01l/l0 ) with temperature for ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites is plotted in Fig. 4. The results in this figure indicate that for each of the composites, the magnitude of \\x01l/l0 increases with temperature up to about 1350  C. The variations of thermal strains, (\\x01l/l0 ) with temperature can be expressed by a third-order polynomial for each of the composites between 200  C and 1350  C. The estimated polynomial equations for the investigated composites  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  565  Table 3 Comparison of the coefficients of thermal expansion (CTE) for ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites at different temperatures. The CTE values obtained for the hot-pressed ZrB2 -SiC (ZS) composites are shown for comparison [39].  Temperature Range (  C)  Coefficient of Thermal Expansion (X 10−6 /K)  20-300 20-400 20-500 20-600 20-700 20-800 20-900 20-1000 20-1100 20-1200 20-1300  ZSBC-10  ZSBC-20  ZSBC-30  ZSBC-40  ZS  3.57 4.63 5.30 5.74 6.06 6.27 6.38 6.53 6.62 6.74 6.70  3.94 4.77 5.35 5.69 5.96 6.14 6.30 6.33 6.44 6.52 6.60  3.79 4.74 5.35 5.75 6.06 6.25 6.3 6.18 6.18 6.33 6.54  2.06 3.44 4.27 4.82 5.21 5.46 5.61 5.72 5.82 5.84 5.78  6.48 7.11 7.48 7.66 7.72 7.78 7.8 7.85 7.9 7.9  having R2 values > 0.9999 are shown as follows:  ((cid:4)l)/l0 (ZSBC-10) = -1.20 ∗ 10 −3 + 7.06 ∗ 10 + 1.17 ∗ 10 −9 ∗T 2 − 6.11 ∗ 10 −13 ∗T3  −6 ∗T  ((cid:4)l)/l0 (ZSBC-20) = -1.13 ∗ 10 + 7.57 ∗ 10 −6 ∗T − 3.99 ∗ 10 −10 ∗T 2 + 1.90 ∗ 10  −3  −13 ∗T3  ((cid:4)l)/l0 (ZSBC-30) = -1.69 ∗ 10 + 10.27 ∗ 10 −6 ∗T − 3.97 ∗ 10 −9 ∗T 2 + 1.51 ∗ 10  −3  −12 ∗T3  ((cid:4)l)/l0 (ZSBC-40) = -1.41 ∗ 10 + 6.13 ∗ 10 −6 ∗T + 2.28 ∗ 10 −9 ∗T 2 − 1.37 ∗ 10  −3  −12 ∗T3  (8)  (9)  (10)  (11)  The coefficients of thermal expansion (CTE) have been calculated for ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites from room temperature to 1350  C using Eq. (4), and the estimated values are compiled in Table 3. The CTE values for the hot-pressed ZrB2 SiC (ZS) composite are also shown in this table for the purpose of comparison [39]. The results in Table 3 show that the CTE of all the composites increases with increasing temperature. Further, the CTE values decrease with increasing SiC content, and the lowest value of CTE is observed for the ZSBC-40 composite. The CTE of SiC is lower than that of ZrB2 , and hence the CTEs of the composites containing more amount of SiC are naturally expected to be lower. Lower matrix grain size of the composites with higher SiC content could also be responsible for their lower CTE values. Larger interfacial area has been found to lower the CTE in case of the TiC particulatereinforced Al matrix composites [32]. The results in Table 4 also show increase in CTE values with increasing temperature range for all the composites, which may be attributed to the transition from residual compressive to tensile stresses acting on the SiC particles [26]. The tensile stress is usually observed at higher temperatures due to greater thermal expansion of ZrB2 compared to that of SiC. Interestingly, comparison of the results in Table 3 show that (i) CTE values obtained within a given temperature range for the pressureless sintered ZSBC-20 are consistently less than those of the hot-pressed ZrB2 -SiC (ZS) composite in spite of finer grain size in the latter material; and (ii) the increase in CTE values with increase  in the temperature range from 20-300  C to 20-1200  C is found to be 258% for the ZSBC-20 in contrast to only 21.9% for the ZS. This observation is suggestive of the role of the interfacial segregation of W-rich phase. The presence of such segregation probably restricts the relaxation at the ZrB2 -SiC interfaces in response to their differential expansion at lower temperatures. However, the increased thermal expansion of the metallic interfacial locations, as well as its ability to plastically deform to accommodate the strain caused by thermal expansion mismatch between ZrB2 and SiC at higher temperatures appears to have contributed to such sharp increase in the CTE of the pressureless sintered composites. For an isotropic two-phase composite, the upper and lower Hashin-Shtrikman bounds of CTE (␣) can be calculated from the following expressions [50]:  (␣eff )u = ␣1 − f2 (␣1 − ␣2 )  (␣eff )l = ␣2 − f1 (␣2 − ␣1 )  2 (3K1 + 4G 1 )  K K1 (3K2 + 4G1 ) + 4f2G 1 (K2 − K1 ) K1 (3K1 + 4G1 ) K2 (3K1 + 4G2 ) + 4f2G1 (K1 − K2 )  (12)  (13)  where (␣eff )u and (␣eff )l are the upper bound and lower bounds, respectively, of CTE of the given composite, whereas ␣, K, G and f are the CTE, bulk modulus, shear modulus and volume fraction, respectively. The subscripts 1 and 2 in Eqs. (12) and (13) represent the matrix and reinforcement phases, respectively. The values of K and G for ZrB2 and SiC used for the calculations have been collected from the literature [49,51]. Further, assuming that a spherical reinforcement is wetted by a uniform layer of matrix, the CTE can be calculated on the basis of Kerner’s model [52,53] using the expression:  ␣c = ¯␣ − f2 (1 − f2 ) (␣2 − ␣1 )  K2 − K1 (1 − f2 ) + K1 f2 + f2 K2 + ((3K1 K2 )/(4G1 ))  (14)  ¯˛ = (1−f2 )␣1 + f2␣2 where (as predicted by the rule of mixtures), while the subscript, c in ␣c stands for the composite. For comparison with the experimentally obtained results, the values of CTE for the investigated composites in the temperature range 20-1000  C have been estimated using various theoretical models including Hashin-Shtrikman relations, Kerner’s model and rule of mixtures by considering the CTEs of ZrB2 (7.17 × 10−6 /K) [31] and SiC (5 × 10−6 /K) [49], and the results are shown in Table 4. For calculations based on the rule of mixtures in this table, the CTE values of B4C (5 × 10−6 /K) [54] and W (3.8 × 10−6 /K) [55] have been also considered, besides those of ZrB2 and SiC. The expressions for CTE derived on the basis of aforementioned models have been already discussed in an earlier report by the authors on the hot-pressed composites [39]. On comparison of the experimental results with theoretical predictions as shown in Table 4, the obtained data are found to be well within the Hashin-Shtrikman limits. The experimentally obtained CTE of ZSBC-40 composite is found to be just a little above (1.02 times) the lower limit, whereas the values found for the composites having ≤30 vol.% SiC particles appear to be greater by 8 to 11%. The ratios of the experimentally determined CTEs for ZSBC-10, ZSBC-20 and ZSBC-30 composites to the upper bound value obtained from the Hashin-Shtrikman relations is 0.94-0.95, whereas it is calculated as 0.85 in case of the ZSBC-40 composite. Using the Kerner’s model, the CTE values are found to be ≈0.94 for the composites with ≤30 vol.% SiC particles, whereas it is calculated as ≈ 0.91 for the ZSBC-40 composite. In a similar manner, the ratios of the experimental results to the values obtained by the rule of mixtures are found to be in the range of ≈0.96-0.97 for the composites with ≤30 vol.% SiC particles, whereas it is calculated as ≈0.94 for the ZSBC-40 composite. Based on the aforementioned comparison of experimental data and results of theoretical predictions, better agreement is found for the investigated composites with ≤30 vol.% SiC compared to that for the composite having 40 vol.% SiC particles.  \\x0c', '566  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  Table 4 Theoretical predictions of coefficients of thermal expansion (CTE) of ZSBC-10, ZSBC-20, ZSBC-30 and ZSBC-40 composites for the temperature range of 20-1000   C along with experimentally obtained results. The ratios of experimentally obtained results to the corresponding predicted values for different composites are shown in parentheses.  Composites  Coefficient of Thermal Expansion (X 10−6 /K)  Present study  Hashin-Shtrikman bounds  Kerner’s model  Rule of mixtures  ZSBC-10 ZSBC-20 ZSBC-30 ZSBC-40  6.53 6.33 6.18 5.72  Lower  5.9 (1.11) 5.8 (1.09) 5.7 (1.08) 5.6 (1.02)  Upper  7.1 (0.92) 7.0 (0.90) 6.9 (0.89) 6.8 (0.84)  6.96 (0.94) 6.75 (0.94) 6.54 (0.94) 6.32 (0.91)  6.77 (0.96) 6.58 (0.96) 6.34 (0.97) 6.10 (0.94)  Furthermore, the predictions based on rule of mixtures are closer to the experimentally obtained values, in comparison to those based on other theoretical models. This observation can be attributed to consideration of the CTEs of B4C and W only in case of the rule of mixtures calculations. The Hashin-Shtrikman bounds and Kerner’s models are basically meant for dual phase materials, and therefore the CTEs of B4C and W could not be considered. It may be noted that ZSBC-40 composite has exhibited the highest amount of interfacial segregation of W-rich phase, which possesses a CTE value different from that of either ZrB2 or SiC. Increased net interfacial area due to the presence of W-enriched layer in the ZSBC-40 composite can be considered to be responsible for not only lowering the CTE, but also greater difference of the experimentally obtained results with theoretical predictions. Comparatively lower value of the experimentally obtained CTE compared to that expected theoretically for the ZSBC-40 composite could also be due to slippage of the relatively weaker ZrB2 -SiC interfaces caused by differences in the amounts of thermal expansion of the constituent phases during heating and cooling cycles.  3.5.  Thermal shock resistance  A comparative assessment of thermal shock resistance of the investigated composites has been made by evaluation of damage caused through thermal cycling by soaking at 800  C, 1000  C or 1200  C followed by quenching in water. The changes in hardness and indentation crack length of each of these composites due to thermal cycling are shown in Fig. 5(a) and (b), respectively. It is evident from Fig. 5 that (i) hardness of each composite decreases with an increasing temperature differential (\\x01T); (ii) net decrease in hardness observed for the ZSBC-20 composite is less than that for the other two composites; and (iii) the indentation crack length of each composite increases with an increasing temperature differential (\\x01T). In other words, the damage due to thermal cycling scales with \\x01T, and is found to be the lowest in case of the ZSBC-20 composite. The SEM micrographs depicting the surface morphologies of the pressure-less sintered ZrB2 -SiC composite specimens, which were exposed at 1000  C for 10 min and then water-quenched, are shown in Fig. 6. Analyses of chemical compositions have shown the formation of a discontinuous scale of ZrO2 . Damage is caused during rapid quenching after exposure at temperatures in the range of 800-1200  C probably due to residual stresses caused by CTE mismatch at the interfaces between ZrB2 and SiC (\\x01CTE ≈ 2.17 × 10−6 K−1 ), as well as due to temperature gradient within the sample. The thermal stress caused by the temperature gradient (\\x01T) is given by the relation [56]: \\u2434th = E␣.\\x01T/(1-(cid:8)) the Young(cid:6) s modulus, ␣ is the CTE, \\x01T is where E is the temperature difference, and (cid:8) is the Poisson(cid:6) s ratio. Therefore, it is obvious that the stress due to thermal gradient within a sample is minimized with decreasing CTE and increasing thermal conductivity. Moreover, the CTE of ZrO2 scale (≈10.3 × 10−6 K−1 ) is much  (15)  higher than that of the composite, whereas its thermal conductivity (≈2.0 W/mK) is lower. Such a large difference of thermal properties of the composite with the ZrO2 scale formed on the surface during high temperature exposure could also contribute to thermal stresses causing spallation and enhanced damage in the investigated composites. The results related to the thermal properties of the composites show that the CTE of ZSBC-20 composite is less than that of ZSBC-10, whereas the thermal conductivity of the former composite is much higher. Furthermore, the CTEs of ZSBC-20 and ZSBC-30 are very close to each other, whereas thermal conductivity of the former composite is much higher. Therefore, superior thermal shock resistance of the ZSBC-20 composite can be attributed partly to its attractive thermal properties. The indentation crack-path in the ZSBC-30 composite as depicted in Fig. 6(d) appears to be along the matrix grain boundaries or ZrB2 -SiC interfaces, indicating that the differential thermal expansion of the W-rich interfacial phase or oxidation of W followed by vaporization of WO3 could play a significant role in damage during thermal cycles particularly on exposure at temperatures ≥1000  C.  3.6. Ablation behavior  3.6.1.  Temperature histories  The variations of front and back surface temperatures with time for ZSBC − 10, ZSBC − 20, and ZSBC − 30 composites are shown in Fig. 7(a) and (b), respectively. The results in Fig. 7(a) show that during the initial stage of heating, the front surface temperature increases sharply for all the composites, and subsequently to 2180  C for ZSBC − 10, it reaches a steady state value equal 2250  C for ZSBC − 20, and 2190  C for ZSBC − 30. In contrast, it is obvious from Fig. 7(b) that the rate of temperature rise at the back surface of the investigated composites is considerably lower than that observed at the front surface. During the initial stage (in till 900 − 990  C, and then the rate of temperature rise decreases. the range of 60-110 s), the back surface temperature rises rapidly This type of variation in temperature can be attributed to decrease in both thermal conductivity with increasing temperature, as well as reduction in the temperature gradient between front and back surfaces. The scientific basis for correlation of the thermal gradient as shown in Fig. 7 with both thermophysical properties and the oxide scale constituents formed during the high temperature exposure has been discussed further in Section 3.6.5.  3.6.2. Morphology and composition of exposed surfaces  Photographs depicting the surfaces of ZrB2 -SiC ceramic composites before and after the ablation tests are shown in Fig. 8, which indicate that an oxide scale is formed on the specimen surfaces exposed to the oxy-acetylene flame. Careful observation of the photographs of specimen surfaces (Fig. 8) also reveals that the oxide scale formed on the ZSBC-20 composite is denser and more adherent compared to that formed on ZSBC-10 or ZSBC-30 composite. Cracks along with evidence of spallation are observed in the oxide scale formed on ZSBC-10 or ZSBC-30 composite, as shown in Fig. 8.  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  567  Fig. 5. Bar charts showing changes in (a) hardness and (b) indentation crack lengths of the ZrB2 -SiC composites due to thermal cycling through exposure at 800, 1000 or 1200   C followed by quenching in water.  Fig. 6. SEM micrographs depicting the oxide scale formed on surfaces of the composites exposed at 1000   C in air and then water-quenched: (a) ZSBC-10, (b) ZSBC-20, and (c) ZSBC-30 composites; as well as (d) indentation crack on the ZSBC-30 composite.  Typical XRD patterns from the oxide scales formed on ZSBC−10, ZSBC−20 and ZSBC−30 composites depict the peaks of monoclinic ZrO2 and ZrSiO4 (as shown in Fig. 9). Pairs of low and high magnification SEM (SE) images, shown in Fig. 10(a) through (f), depict the typical morphologies of the oxide scales formed on ZSBC-10, ZSBC-20, and ZSBC-30 composites. Chemical compositions of these scales have been studied by EDX analyses, and a typical spectrum is shown in Fig. 10(g). The presence of the peaks of Zr, Si and O in the spectrum suggests that the oxides of Zr and Si are present in the oxide scale, which is consistent with the results of XRD analyses. Micro-porosities are observed from the examination of the surface morphology of the oxide scales, some of which are linked by networks of cracks, while the rest are found to be isolated and distributed randomly. It is confirmed from the qualitative comparison of the SEM images that: (i) the oxide scale of the ZSBC-20 composite has both the finest and the lowest density of cracks among the presently investigated composites, whereas (ii) the densest network of cracks and the widest discontinuities are visible in the oxide scale of the ZSBC-30 composite. A typical SEM (SE) image depicting the cross-section of the oxide scale and corresponding EDX elemental maps of Zr, Si and O for ZSBC-20 composite after ablation for 600 s are shown in  Fig. 11 (a) through (d), respectively. Examination of this figure indicates the presence of three distinct layers. The outer layer is mainly composed of Zr-Si-O reached phase. It is expected that a silica-rich glassy phase is embedded within the pores surrounding the ZrO2 grains. The formation of Zr-Si-O layer effectively hinders the ingress of oxygen, and thereby prevents further oxidation. The intermediate layer appears to be depleted in SiC. This layer is expected to form by active oxidation of SiC. The innermost layer comprising the part of microstructure beneath the SiC depleted region is found to be unaltered, as shown in Fig. 11(a-d). The thicknesses of the oxide scales (excluding the SiC-depleted layer) of ZSBC-10, ZSBC-20 and ZSBC-30 composites have been measured as 196 ± 13 \\u242em, 177 ± 32 \\u242em and 185 ± 15 \\u242em, respectively. Compared to the oxide scales formed on ZSBC-10 and ZSBC-30 composites, that formed on the ZSBC-20 composite has been found to be thinner for identical test condition, which confirms that the amount of degradation of the ZSBC-20 is relatively less compared to that of either ZSBC-10 or ZSBC-30 composites. In a material with lower thermal conductivity, temperature gradient would be less, and therefore the temperature at a given depth beneath the oxide scale would be higher, which in turn would increase the thickness of the oxide scale, as has been observed in  \\x0c', '568  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  Fig. 9. X-Ray diffraction patterns obtained from the oxide scales formed on ZSBC−10, ZSBC−20 and ZSBC−30 composites after subjecting to ablation for 600 s.  case of ZSBC-10 or ZSBC-30 composites. As shown in Fig. 8, the outer layer of the oxide scale in ZSBC-10 or ZSBC-30 composite is swollen and loose, which appears to promote spallation during cooling. Higher thickness of the oxide-scale formed on ZSBC-10 or ZSBC-30 composite is also expected to be responsible for its higher residual stress as well as its loose nature. The residual stress is generated in the oxide scale during cooling due to the mismatch between the CTEs of oxide scale and the composite substrate beneath, phase transformation in ZrO2 , as well as the pressure exerted by escape of CO, B2O3 and WO3 . The stress is relaxed by cracking and spallation of parts of the oxide scale. Intuitively, a strongly adherent oxide scale formed at the surface of the ZSBC-20 composite has a lower thermal conductivity and would act as a thermal barrier coating, which may explain the reason for less damage during exposure at elevated temperatures.  3.6.3.  Effect of ablation on mass gain and elastic modulus  Bar-charts depicting the changes in mass and Young’s modulus during high temperature exposure of the ZrB2 -SiC composites are shown in Fig. 12(a) and (b), respectively. The results depicted of the ZSBC − 10 composite, and the lowest for the ZSBC-20; and in these figures indicate that (i) mass gain is the highest in case (ii) the Young’s modulus recorded after exposure for ablation is lower than that of pre-ablation tested samples by 21%, 10% and 20% for ZSBC-10, ZSBC-20 and ZSBC-30 composites, respectively. The decrease in Young(cid:6) s modulus can be attributed to generation of internal flaws caused by thermal stresses owing to temperature gradient in the composite samples as well as their exposure at 2200  C. The formation of oxide scale with cracks and porosities on the surface along with partial spallation (as is evident from Figs. 8 and 10) has also contributed to reduction of Young(cid:6) s modulus. Therefore, the observed trend regarding decrease in Young(cid:6) s modulus confirms that the degradation due to ablation is the least for the ZSBC-20 composite among the investigated pressureless sintered composites.  3.6.4. Ablation mechanism  Analyses of the oxide scale microstructures as well as changes in mass and Young(cid:6) s modulus indicate that the ZSBC-20 composite exhibits less environmental degradation by ablation than either ZSBC-10 or ZSBC-30 composite. During re-entry, the process of ablation involves high temperature oxidation, erosion caused by the high speed air flow, as well as vaporization of volatile oxidation products. In the present experimental study, high speed air flow  Fig. 7. Plots showing the variations of (a) front and (b) back surface temperatures (  C) of ZrB2 -SiC composites with duration of exposure to neutral oxy-acetylene flame.  Fig. 8. Photographs of the surfaces of ZSBC-10, ZSBC-20 and ZSBC-30 composites (a) before ablation, and (b) after ablation.  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  569  Fig. 10. SEM (SE) images of: (a and b) ZSBC−10, (c and d) ZSBC−20, (e and f) ZSBC−30 composites subjected to ablation test for 10 min; and (g) a typical EDX spectrum from the oxide scale.  was missing, and therefore mass loss by spallation is not observed. It has been shown that formation of a protective and adherent oxide scale at the surface provide adequate configurational stability to the ZrB2 -SiC composite, when subjected to the simultaneous application of mechanical load and oxidation [19,57]. Therefore, it is intuitive that the ZSBC-20 composite showing the formation of a stable oxide scale would be most resistant to ablation. The results of this study differ from the observation of Han et al. [58], who have reported a mass loss of ≈−0.23 mg/s during exposure of the ZrB2 -20 vol.% SiC composite to oxy-acetylene torch at 2200  C. In this study, the result is different probably because the mass-gain  by oxidation significantly exceeds the amount of mass-loss during the high temperature exposure. During high temperature exposure, reactions are expected to occur:  following oxidation  the  (2/5)ZrB2 (s)) + O2 (g) → (5/2)ZrO2 (s) + (½)B2O3 (l)  B2O3 (l) → B2O3 (g)  (2/3)SiC(s) + O2 (g) → (2/3)SiO2 (l)) + (2/3)CO(g)  ZrO2 (s)) + SiO2 (l) → ZrSiO4 (s)  (16)  (17)  (18)  (19)  \\x0c', '570  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  Fig. 11. Cross-section of the oxide scale formed on ZSBC-20 by exposure for 10 min during ablation test: (a) SEM (SE) image, as well as EDX X-Ray maps of: (b) Zr, (c) Si and (d) O.  SiO2 (l) → SiO(g)) + (½)O2 (g)  (2/3)W(s)) + O2 (g) → (2/3)WO3 (g)  (22)  (23)  The major constituents of the oxide scale have been found to be ZrO2 , SiO2 and ZrSiO4 , as shown in Figs. 9-11. During the process of oxidation, gases and volatile products like BO, B2O3 , CO and CO2 are evolved, which escape by leaving behind cracks and pore-channels in the oxide scale, as is evident from Reactions (16) to (18) as well as Reactions (20) and (21). During the ablation tests, the surface temperatures of the composites reach 2200  C. At this temperature, the equilibrium partial pressures of oxygen for oxidation of ZrB2 and SiC at 2200  C have been reported as 10−10 and 10−9 atm, respectively [58], which suggests that the latter phase is more prone to oxidation. The effective protection from SiC is weakened at this temperature due to its active oxidation according to Reactions (20) and (21), followed by formation and escape of volatile SiO (Reactions (20) and (21)) [59-61]. The stability of the oxide scale depends on volatilization and decomposition of the oxidation products with high vapor pressures. The vapor pressures of B2O3 , SiO2 and ZrO2 at 2200  C are 2.76 atm, 1.77 × 10−4 atm and 2.47 × 10−7 atm, respectively [61], indicating order of decreasing volatility of the above-mentioned oxidation products. At high temperatures, SiO2 (l) either decomposes to form SiO (g) through Reaction (22), or reacts with CO (g) formed by oxidation of SiC through Reactions (18)-(21) to form SiO (g). Further, B2O3 also vaporizes rapidly on heating beyond 1100  C. The escape of volatile gaseous products causes degradation of the initially formed protective glassy scale by leading to its rupture and allowing ingress of oxygen from the oxide − air interface to the composite substrate underneath, which in turn enhances the oxidation kinetics. Among the oxidation products, ZrO2 has the lowest vapor pressure, and therefore it is stable. However, the ZrO2 scale is porous and develops cracks due to phase transformations during heating and cooling cycles, which is evident on examination of Fig. 10. On comparison of the XRD patterns (Fig. 11), it is observed that the relative intensities of the XRD peaks of ZrSiO4 as compared to those of other constituent phases of the oxide scale are higher for the ZSBC-20 composite than that for other investigated composites. On examination of ZrO2 -SiO2 phase diagram, it is observed that  Fig. 12. Bar charts showing changes in (a) mass and (b) Young(cid:6) s modulus in the composites with varying SiC content after 10 min of ablation tests at 2200   C.  SiC(s)) + O2 (g) → SiO(g)) + CO(g)  SiC(s)) + 2SiO2 (l) → 3SiO(g) + CO(g)  (20)  (21)  \\x0c', 'M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  571  a liquid phase co-exists with cubic ZrO2 at 2200  C towards the ZrO2 -rich side, and the amount of liquid increases with increasing SiO2 content [62]. During cooling, this liquid solidifies into ZrO2 + crystoballite at 1960  C, which is converted at 1949  C to mixtures of ZrSiO4 + ZrO2 or ZrSiO4 + SiO2 , depending on the SiO2 content of the oxide scale. It is intuitive that formation of ZrSiO4 occurs in the oxide scale during heating by Reaction (15) between ZrO2 and SiO2 . However, ZrSiO4 is expected to decompose into ZrO2 and crystalloballite with increase in temperature beyond 1949  C, and this is followed by formation of liquid phase along with ZrO2 on further heating. Formation of this liquid is expected to aide in closure of porosities along with sintering of the oxide scale, and in this manner create a barrier for retarding the diffusion of oxygen anions. The ZSBC-10 composite having the lowest SiC content would also form a relatively lower amount of SiO2 in its oxide scale, and therefore amount of liquid phase formation at high temperature is expected to be less. On the other hand, the ZSBC-30 composite with higher SiC content may show higher degree of active oxidation. However, the ZSBC-20 composite can be considered to have nearly optimum ratio of ZrO2 to SiO2 in its oxide scale, such that the latter oxide at the surface is consumed to a larger extent for formation of ZrSiO4 , such that a sufficiently large amount of liquid forms during high temperature exposure. This mechanism could explain the reason for higher smoothness as well as relatively smaller size of pores found in the oxide scale of the ZSBC-20 composite compared to that of either ZSBC-10 or ZSBC-30 composite (Fig. 10). A small fraction of mass-loss could be due to formation and vaporization of WO3 , as the exposure temperature is much higher than that of its sublimation [63]. Considering that the W-content of the ZSBC-30 composite is only 2.8 wt.% (≈1.2 at.%), the damage due to vaporization of WO3 (Reaction (23)) is not expected to be significant, and certainly its role would be even less important in case of ZSBC-10 and ZSBC-20 composites with W concentration of 0.52 and 0.91 at.%, respectively.  3.6.5. Relation of thermal gradient with thermal properties and oxide scale  The front surface temperature of a material depends on the emissivity, thermal conductivity and specific heat capacity. The insignificant difference in the temperature profiles as shown in Fig. 7, in spite of the differences in thermal diffusivities of the investigated composites could be due to the following reasons: (i) irrespective of the SiC volume fraction, the top surface of each sample shows the formation of ZrO2 + liquid followed by a SiC-depleted region underneath, and (ii) the precision with which temperature could be recorded manually was limited, as the initial rise in the temperature to 2000  C took only 60-110 s. The ZrO2 -rich scale, which acts as a thermal barrier coating during exposure at 2200  C, has more or less identical emissivity, until it becomes thick enough and undergoes spallation. Higher surface temperatures of the ZSBC20 composite after reaching the steady-state condition as shown in Fig. 3 may be attributed to its higher emissivity compared to that of ZSBC-10 or ZSBC-30, which have shown the presence of larger cracks and porosities in the oxide scales. However, further experiments are required to measure the emissivity and correlate it with the front surface temperature. It is also observed that the rise in the back surface temperature is a little faster for the ZSBC-20 composite with higher thermal conductivity, compared to either ZSBC-10 or ZSBC-30.  4. Conclusions  In the present study, the effect of SiC volume fraction as well as segregation of W-rich phase at the interfaces in the pressure less sintered ZrB2 -SiC composites on electrical resistivity, thermal conductivity, coefficient of thermal expansion, thermal shock and ablation resistance have been examined. The electrical resistivity of the investigated pressureless sintered ZrB2 -SiC composites increases more than four times with both increasing volume fraction of SiC, and with increasing area of matrix grain boundaries. Values of average interfacial electrical resistance increase by more than 15% with volume fraction of SiC in the composite, because of its lower electrical conductivity than that of the ZrB2 matrix. The thermal conductivity of the ZSBC-20 at various temperatures is found to be higher by 3-24% compared to that of the other investigated composites. The thermal conductivity appears to depend on relative density and SiC content, as well as on interfacial resistance beyond a critical SiC volume fraction of 20%. The contribution of phonon-scattering to thermal diffusivity increases with increasing SiC content, and therefore the thermal conductivity of the ZSBC-30 is lower by 1-8% compared to that of ZSBC-10 temperatures > 900  C. ues in the temperature range of 20-1000  C decrease by 3-12% at It has been observed that the CTE valwith increase in the SiC content from 10 to 40 vol.%. For all the investigated composites, the CTE values increase with increasing temperature, due to increase in relaxation and slippage at interfaces. Due to the interfacial segregation of W, Fe, and Co in the pressureless sintered ZSBC-20 composite, its electrical resistivity has been found to be less by 2.8 times compared to that of the hotpressed composite with clean ZrB2 -SiC interfaces, whereas thermal conductivity at 1200  C and CTE in the range of 20-1200  C are less by 29.1% and 21.1%, respectively. Compared to the ZSBC-20 composite, ZSBC-10 and ZSBC-30 have exhibited reduction in hardness values by 2.4 and 1.4 times, respectively due to damage on exposure to thermal shock by quenching after soaking at 1200  C. This observation is attributed to desirable combination of thermal conductivity and expansion coefficient in case of the ZSBC-20 composite. Further, the ZSBC-20 has exhibited superior resistance to damage by ablation and thermal shock due to exposure at 2200  C followed by air-cooling, as compared to those of either ZSBC-10 or ZSBC-30 composite. This inference is drawn considering that the mass-gain recorded for the ZSBC-20 composite due to oxidation is less by 5-10%, whereas the decrease in Young(cid:6) s modulus is less than half compared to that of ZSBC-10 or ZSBC-30 composite. Relatively less damage in case of the ZSBC-20 composite is attributed to both superior thermal properties and optimum amounts of ZrO2 and SiO2 in the oxide scale to form ZrSiO4 , which decomposes on further heating to form mixture ZrO2 + liquid offering protection against oxidation. In the ZSBC-30 composite, active oxidation of SiC could create the potential path for the inward transport of oxygen during high temperature exposure.  Acknowledgements  The financial support from the Defence Research and Development Organization, New Delhi is gratefully acknowledged. The authors also express their sincere gratitude to Mr. Mithun Das, Mr. Santu Mudliyar, and Mr. Ronadhir Bosu, technicians at the Central Research Facility, IIT Kharagpur for providing assistance in characterization.  References  [1] F. Monteverde, A. Bellosi, Microstructure and properties of an HfB2 -SiC composite for ultra high temperature applications, Adv. Eng. Mater. 6 (2004) 331-336. [2] P. Kolodziej, Aerothermal performance constraints for hypervelocity small radius unswept leading edges and nose tips, NASA Technical Memorandum (1997) 1122047.  \\x0c', '572  M. Mallik et al. / Journal of the European Ceramic Society 37 (2017) 559-572  [3] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra-high temperature ceramics for aero-propulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767. [4] A.L. Chamberlain, W.G. 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  "_id": 52,
  "PDF": "Effect of SiC Content on the Ablation and Oxidation Behavior of ZrB2-Based Ultra High Temperature Ceramic Composites.pdf",
  "Text": "['Materials 2013, 6, 1730-1744; doi:10.3390/ma6051730   Article   OPEN ACCESS  materials   www.mdpi.com/journal/materials   ISSN 1996-1944   Effect of SiC Content on the Ablation and Oxidation Behavior  of ZrB2-Based Ultra High Temperature Ceramic Composites   Ping Hu *, Kaixuan Gui, Yang Yang, Shun Dong and Xinghong Zhang   Science and Technology on Advanced Composites in Special Environment Laboratory, Harbin  Institute of Technology, Harbin 150001, China; E-Mails: gkx89@sina.com (K.G.);  18746026280@163.com (Y.Y.); 18746046862@163.com (S.D.); zhangxh@hit.edu.cn (X.Z.)   * Author to whom correspondence should be addressed; E-Mail: huping@hit.edu.cn;  Tel.: +86-0451-8640-3016; Fax: +86-0451-8640-3016.    Received: 9 February 2013; in revised form: 28 March 2013 / Accepted: 29 March 2013 /   Published: 29 April 2013   Abstract: The ablation and oxidation of ZrB2-based ultra high  temperature ceramic  (UHTC) composites containing 10%, 15% and 30% v/v SiC were tested under different  heat fluxes in a high frequency plasma wind tunnel. Performance was significantly affected  by the surface temperature, which was strongly dependent on the composition. Composites  containing 10% SiC showed the highest surface temperature (>2300 °C) and underwent a  marked degradation under both conditions. In contrast, composites with 30% SiC exhibited  the lowest surface temperature (<2000 °C) and demonstrated excellent ablation resistance.  The surface temperature of UHTCs in aerothermal testing was closely associated with the  dynamic evolution of the surface and bulk oxide properties, especially for the change in  chemical composition on  the exposed surface, which was strongly dependent on  the  material composition and testing parameters (i.e., heat flux, enthalpy, pressure and test  time), and in turn affected its oxidation performance.   Keywords: ZrB2; SiC; ultra high temperature ceramics (UHTCs); ablation; oxidation    1. Introduction   Refractory metal borides such as zirconium diboride (ZrB2) and hafnium diboride (HfB2) have been  commonly referred to as ultra high temperature ceramics (UHTCs), for their extremely high melting  temperatures (around 3300 and 3500 K respectively) [1]. ZrB2 and HfB2 based UHTC composites         \\x0c', 'Materials 2013, 6   1731   represent a class of promising materials for use in extreme applications such as sharp leading edge and  control surface components on hypersonic vehicles, because of their high melting point, retained  strength at elevated temperatures, relatively good oxidation resistance, and dimensional stability in  hypersonic flight conditions [2-8]. Zirconium diboride has some advantages over hafnium diboride as  an aerospace material because it is lighter and less expensive while having comparable levels of  oxidation resistance [4,7-12].  ZrB2-SiC composites are currently considered  the baseline ultra high  temperature ceramic  composites. Indeed, varying the starting composition by changing the SiC content has given added  flexibility in optimizing specific microstructure designs. Adjusting the SiC content in ZrB2 matrix, for  instance, has proven beneficial  for oxidation and ablation  resistance  [3,7,9-14]. Specifically,  compositions containing from 10% to 30% (by volume) SiC have generally been found to be optimal  in this regard. The oxidation resistance of ZrB2-SiC of various compositions under the same oxidizing  condition has been studied [7,14]. However, these have focused primarily on static or flowing air  studies using furnace testing, and few studies have evaluated their oxidation behavior in simulated  hypersonic flight environments [9-13]. Oxidation testing results involving both furnace oxidation  testing and plasma wind  tunnel  testing (or arc-jet  testing) showed  that UHTCs have a similar  microstructure of the oxide scale and oxidation resistance at a comparable level of the sample surface  temperature [7-16]. A high-temperature thermal protection system (TPS) intended for the leading edge  and control surface components of a hypersonic vehicle will likely encounter partially dissociated air  in chemical non-equilibrium with the TPS surface. This will cause a different surface temperature for  samples with different compositions under the same testing conditions [17]. Unfortunately, data  regarding the response of material composition on the surface temperature of UHTCs is not available.  Reducing the surface temperature of the materials by changing composition may be a good way to  greatly improve the material performance in hypersonic conditions, thus enabling hypersonic vehicles  to operate under higher heat flux conditions. How to decrease the surface temperature is an urgent  issue for aerospace applications, and much more attention should be paid to the thermal response of  UHTC composites to the service environment in the future.  The purpose of this paper is to investigate the effect of SiC content on the performance of ZrB2  based ultra high temperature ceramic composites under different heat flux using high frequency plasma  wind tunnel. The ablation resistance, oxidation behavior and thermal response of these composites  is discussed.   2. Experimental Procedure   The samples for ablation tests were fabricated from commercial ZrB2 (Northwest Institute for   Non-ferrous Metal Research, China) and SiC (Weifang Kaihua Micro-powder Co. Ltd., China)  powders. The ZrB2 and SiC powders had the same purity of 99.5% and their mean particle sizes were  2 μm and 0.5 μm respectively. The powder mixtures of ZrB2 with 10% SiC v/v (ZS10), ZrB2 with   15% v/v SiC (ZS15), and ZrB2 with 30% v/v SiC (ZS30) were ball milled in ethanol for 8 h with hard  milling media, and dried in a rotating evaporator. Milled powder was then uniaxially hot pressed in a  boron nitride coated graphite die at 1950 °C for 60 min under vacuum and 30 MPa of applied       \\x0c', 'Materials 2013, 6   1732   pressure. All the samples have a nearly full density and the effect of porosity on the oxidation behavior  can be ignored.  The samples with a Φ20 mm × 30 mm cylinder for ablation tests were cut from the billet, and all  surfaces were diamond polished to a 1 μm finish then exposed to sustained high enthalpy flows using a  high frequency plasma wind tunnel. Coupons were ultrasonically cleaned, successively in detergent,  de-ionized water, acetone and alcohol prior to their exposure. Test specimens were mounted in  graphite holders attached to a water cooled sting arm. The cold wall stagnation point heat flux was  measured using a water cooled copper calorimeter installed flush with the surface of a water cooled  copper probe. The performance of ZrB2-SiC UHTCs in subsonic high enthalpy dissociated airflows  was  investigated. The gas mass flow rate was 3-4 g/s. The other main  testing parameters are  summarized in Table 1. The experiments were carried out with a two-colour Raytek pyrometer  (RAYMR1SCSF, USA), which covers a temperature range of 1000 to 3000 °C. X-ray diffraction  (Rigaku, Dmax-rb, Japan), scanning electron microscopy (FEI Sirion, the Netherland), and energy  dispersive spectroscopy (FEI Sirion, the Netherlands) were used to characterize the phase composition  and microstructure of the surface, and the cross section of the samples after testing.   Table 1. Testing parameters.   Condition   Heat flux (MW/m2)  Enthalpy (MJ/kg)   Pressure (Kpa)   1  2   4.78  3.80   27.9  20.8   18  18   3. Results   ZS10 and ZS15 displayed significant ablation after tests under condition 1, whereas ZS30 exhibited  excellent ablation resistance and configuration stability, as shown in Figure 1. Detailed ablation results  are given in Table 2. The oxide scales were not adherent to the base material for ZS10 or ZS15. Rough  surfaces, along with some holes and crests, were detected on these two composites, besides cracks. In  contrast, a smooth surface was observed for ZS30, and the formed oxide scale was adherent to the base  material. It should be noted that there was a marked difference on the surface temperature between  these composites. ZS10 displayed the highest surface temperature, and the lowest surface temperature  was found in ZS30. In addition, a temperature jump was detected for ZS10, as shown in Figure 2.  Comparing the results of the two conditions, it appears that the surface temperatures are in good  agreement and that the surface temperatures obtained for a sample under condition 2 are lower than  those under the condition 1. ZS30 exhibited the lowest surface temperature in both conditions. In  particular, although the heat flux in condition 1 was higher than that in condition 2, the surface  temperature of ZS30 under condition 1 was still lower than that of both ZS10 and ZS15 under  condition 2.           \\x0c', 'Materials 2013, 6   1733   Table 2. Ablation results of ZrB2-based ultra high temperature ceramic composites.   Materials   Condition   Original  mass (g)   Mass after  ablation (g)  Steady state surface  temperature (°C)   Ablation  time (s)   ZrB2-10 vol % SiC  ZrB2-15 vol % SiC  ZrB2-30 vol % SiC  ZrB2-10 vol % SiC  ZrB2-15 vol % SiC  ZrB2-30 vol % SiC   1  1  1  2  2  2   51.44  53.61  49.81  51.54  53.68  49.76   50.99  53.52  49.78  51.18  53.66  49.77   2450  2320  1960  2330  2250  1900   300  600  600  600  600  600   Figure 1. Photographs of ZrB2 based ultra high temperature ceramic composites during and  after ablation tests under condition 1: (a) ZS10; (b) ZS15; (c) ZS30.   a   b   c   Figure 3 shows the photographs of ZrB2-SiC composites after ablation tests under condition 2.  Spallation was not observed after cooling to room temperature because the heat flux was lower than  that  in condition 1. ZS30 also showed  the best ablation and oxidation resistance among  these  composites under the condition 2, which is consistent with the results under condition 1. The surface  colors of these composites became grayer with increasing SiC content. ZrO2 appeared white, while         \\x0c', 'Materials 2013, 6   1734   silica glass appeared dark at room temperature, meaning that the amount of the silica glass at the  surface increases with increasing SiC content.    Figure 2. Images of ZS10 during ablation tests under condition 1.   Figure 3. Photographs of ZrB2 based ultra high temperature ceramic composites after  ablation tests under condition 2: (a) ZS10; (b) ZS15; (c) ZS30.   a   b  c  The oxide scale consisted primarily of zirconia and showed an oriented growth during ablation  testing of ZS10 under condition 1 (Figure 4). A columnar shape of the scale, containing some voids,  was detected after ablation. Such observations are consistent with the furnace oxidation of these  composites being at temperatures of 1800 °C and above [14-16,18]. The outside oxide scale easily  cracked, or spalled, during cooling to room temperature, owing to the different CTEs between oxide           \\x0c', 'Materials 2013, 6   1735   scale and the base material, as shown in Figure 4. The phase transformation of zirconia, with an  accompanying volume change, can easily  lead  to cracking and spalling upon cooling  to room  temperature. Furthermore, the formation of high pressure gas phases (e.g., B2O3, CO, and SiO)  resulting from the oxidation of ZrB2 and SiC was responsible for the rupture of the oxide scale during  the oxidation process. These are possible mechanisms to explain the scale loss that was observed. High  magnification of the oxide scale showed clear evidence of gas flow, which indicated that gaseous  products were escaping from the material. In addition, a large number of pores were embedded in the  ZrO2 rich oxide scale. Such observations have not been detected in previous studies. Cross-sectional  morphologies of ZS10 after ablation under condition 1 revealed that the top surface layer had flaked  off, as shown in Figure 5. The total pressure in the inner oxide scale is a key factor in the mechanical  stability of the oxide scale. Thermodynamic calculations indicated that at an air temperature of   2066 °C, the total pressure of B2O3 at the ZrB2/ZrO2 interface would reach 1 atm [19]. Thus, the total  pressure of the interior oxide scale was higher than ambient pressure, since the surface temperature  was far higher than 2066 °C and the ambient pressure was lower than 1 atm under the conditions used.  The side surface scale was adherent to the base material because the temperature in this region was  significantly lower than that at top surface. Figure 6 shows SEM images of the substrate in which the  outside scale was lost. The surface was not smooth, and a large number of concave pits were detected.  The growth of oxide scale with a cylindrical shape in the concave pit is obvious from Figure 6c.  Further growth of these cylindrical grains would sinter together with the outside scale and lead to an  increase of the thickness of the outside scale. Pores were commonly detected at the grain boundaries  and the inner grains, as a result of the outflow of the gaseous phase products. This observation is  consistent with the microstructure of the outside scale, as shown in Figure 4. Interestingly, the  microstructure and morphology of the substrate showed a particle shape (Figure 6d) similar to the ZrB2  grains. In addition, a number of cracks were detected in the grains due to volume expansion on the  conversion of ZrB2 to ZrO2. EDS analysis of the substrate layer showed that ZrB2 was not completely  transformed to ZrO2 (not shown). The apparent growth of ZrO2 has not occurred in this region. We  also found that surface of the substrate had an intact structure, but little evidence of fracture was  observed. Such observations  indicate  that  the outside scale was not effectively adherent  to  the  substrate, which was further confirmed by the fact that the formed outside scale flaked off by itself at  room temperature.    Figure 4. SEM micrographs of the oxide scale for ZS10 after ablation under condition 1:  (a) low magnification; (b) high magnification.   a   b  100 μm  20 μm         \\x0c', 'Materials 2013, 6   1736   Figure 5. Cross-sectional morphologies of ZS10 after ablation under condition 1: (a) low  magnification; (b) high magnification of the scale at the side surface.   a   resin  b  500μm  100μm   Figure 6. SEM micrographs of the substrate for ZS10 after ablation under condition 1:   (a) low magnification; (b), (c) and (d) are higher magnifications of (a).   a   c   b  d  300μm  50μm   10μm  5μm   SEM micrograph of the exposed surface of ZS30 after ablation testing under condition 2 revealed a  coherent and compact silica glass coating decorated with aggregated zirconia crystals of various sizes  and shapes, as shown in Figure 7. In contrast, the surface was not smooth and contained a number of  pores on the surface after ablation under condition 1. Moreover, the amount of the silica glass formed on  the surface  layer under condition 1 was much  lower  than  that formed under condition 2. Such  observations are attributed to the different surface temperature as a consequence of the change of the test  condition. Figure 8 shows the cross-sectional morphologies of ZS30 after ablation tests under the two  conditions. The oxide scale formed under condition 2 was more compact than that formed under           \\x0c', 'Materials 2013, 6   1737   condition 1. Moreover, the formed ZrO2 had not apparently changed the initial frame structure of the  ZrB2 after ablation under condition 2. A silica rich outer layer was formed under the present condition, as  shown in Figure 8. External silica scales, such as those observed in the present study, have been shown  to limit the inward diffusion of oxygen, thus enhancing the resistance to oxidation [7,9]. The thickness of  oxide scale for ZS30 ablated under condition 2 for 600 s was only ≈90 μm, which is significantly lower  than that of ZS10 and ZS15 under the same condition. Under harsher conditions, the silica rich glass  layer was not observed in the cross-section, which is consistent with surface microstructure. Silica glass  became unstable and was lost by rapid vaporization in the present case. The elemental mapping by   SEM-EDS throughout the oxide scale after ablation at condition 1 for ZS30 revealed the formation of a  two-layer structure, as can be seen in Figure 9. The silica glass were formed and homogeneously  distributed in outside scale, which effectively plugged pores and sealed ZrO2 boundaries leading to an  enhanced oxidation resistance. A little silica glass was observed in the inner oxide scale, as a result of the  active oxidation of SiC. The ZrO2 skeleton was continuously formed, and no evidence of cracking or  spallation was detected (Figure 9d). ZrO2 provides an oxide skeleton without being blown away by the  gas flow, leading to configurational stability of the material. To ensure good resistance, the protective  condensed-phase oxide scale must be thermochemically and mechanically stable. ZS30 exhibits passive  oxidation with parabolic kinetics under high heat flux, as shown in Figure 10. This indicates that the  formed scale acts as an effective barrier for diffusion of oxygen to the underlying material.   Figure 7. Surface morphologies of ZS30 after ablation under (a) condition 1 and (b) condition 2.   a   b  50μm  50μm   Figure 8. Cross-sectional morphologies of ZS30 after ablation under (a) condition 1 and  (b) condition 2.   a   b  40μm  50μm           \\x0c', 'Materials 2013, 6   1738   Figure 9. Elemental maps for the oxide scale of ZS30 after ablation under condition 1.   a a  c c  b b  EDS Map for O EDS Map for O  50μm 50μm  EDS Map for Si EDS Map for Si  50μm 50μm  d d  EDS Map for B EDS Map for B  50μm 50μm  EDS Map for Zr EDS Map for Zr  50μm 50μm  Figure 10. Thickness of the reacted layer vs. ablation time obtained under condition 1.   )  m m  3   0 1  (  120  r e  y a  l  d  e  t  c  a  e r  e  h  t  f  o  s s  e  n k  c  i  80  40  0  h  T  0  200  400  Ablation time (s)  600                          \\x0c', 'Materials 2013, 6   4. Discussion   1739   The ablation and oxidation of UHTC samples exposed to frequency plasma wind tunnel suggests  that SiC content has a significant impact on the surface temperature and oxidation resistance. This can  be explained by the evolution of the surface composition and microstructure, which is the main reason  for  the  surface  temperature difference. Unfortunately,  the major  factor  that accounts  for  the  temperature difference for UHTCs with different compositions under  the same  test  is not well  understood. While ZrB2-SiC ultra high temperature ceramic composites have been studied extensively  in the space field and have been fully characterized in their base thermo-mechanical properties, little  attention has been paid to their radiative and surface catalytic behavior at ultra high temperatures.  Nonetheless, emissivity and surface catalycity represent key parameters for space re-entry thermal  protection system applications. When a space vehicle flies in a very high Mach number, a very strong  bow shock is formed in front of the body, and the dissociation of air occurs in the shock layer. Some of  the dissociated atomic nitrogen and atomic oxygen recombine on the surface. Since recombination  processes are exothermic (498 kJ released per mole of O2 formed, 945 kJ per mole of N2, and   415 kJ per mole of NO) [20,21] part of the released energy is transferred to the surface of materials  with the consequent increasing of their surface temperature [22-24]. The influence of SiC content on  the surface temperature has not been recognized in the previous literature. Marianne Balat-Pichelin et al.  have studied the influence of the microstructure of the material in the case of silica and alumina, and  indicated that the surface morphology, and particularly roughness, play a very important role in the  atomic oxygen recombination coefficient γ0, which increases with increasing surface roughness [25].  Significant swelling of the oxide scale occurred during the oxidation of low SiC content material, due  to the volume expansion as a result of ZrB2 oxidation. This led to the formation of a rough surface  layer during oxidation, which may have caused the increment of the atomic recombination coefficient,  resulting in an increase of surface temperature. It has been reported that silica glass has very weak  catalytic activity and high radiative efficiency compared with other materials [26,27]. Therefore, the  existence of silica glass on the surface would lead to a decrease in surface temperature. High content of  silica glass was observed on the surface of ZS30 after ablation, which would account for it having the  lowest surface temperature among all the materials.  The different oxidation rate of these materials under the same condition may be another reason for the  surface temperature difference, since the oxidation processes are also exothermic. Surface oxidation  processes that transform ZrB2 and SiC into condensed and volatile oxides are all exothermic, and release  more energy per oxygen atom consumed than catalytic recombination to molecular oxygen [17]. Thus,  the effect of the oxidation on the surface temperature cannot be neglected, and the oxidation rate also  plays an important role in the heat formation that led to an increase of surface temperature. It should be  noted that the difference among the surface temperatures of the UHTC specimens with low and high SiC  content was small during initial heating and until a steady state was reached. Obviously, the oxidation  rate will decrease, and at  least not  increase with  increasing  time, when  the sample reaches  the  equilibrium state, resulting in a reduction of heat formation. From the exothermic point of view, the  oxidation will not lead to the further increase of the surface temperature when the UHTC specimens  achieve a steady state. In fact, a large temperature jump, with a magnitude of 300-500 °C, was found in  UHTCs containing low SiC, resulting in higher surface temperatures. Therefore, the different oxidation       \\x0c', 'Materials 2013, 6   1740   rates of ZrB2-SiC composites are not the main factor for the large surface temperature differences,  because the oxidation rates of these materials were all relatively low before the temperature jump, under  the conditions used in this study. Of course, the heat contribution to the heat flux coming from the  material oxidation should be taken into account where the material has a very high oxidation rate.  Similar phenomena was also detected by Marschall J. et al. who concluded that a temperature jump of  ZrB2-SiC UHTCs was associated with a transition in surface chemistry involving the Si-containing  compounds  (the silica-rich oxide scale and  the silica  former SiC) and  is  likely  related  to  the   passive-to-active oxidation transition observed for other SiC-containing composites in aerothermal test  environments [17]. The total emittance values for the oxides of ZrB2-SiC UHTCs are shown to change  slightly over a large temperature range [9]. So, the change in emittance of the oxide scale is not sufficient  to be the driving force for the large temperature jump, according to thermal radiation equilibrium  assessment. Moreover, the total emittance for systems of similar materials has been shown to change  slightly before and after the temperature jump occurs [17]. Therefore, the change in emittance during the  ablation process may not contribute to the large surface temperature difference in the present study.   According to the discussion above, the possible main reasons for the large surface temperature  difference are the increase of the chemical heating contribution, the decrease of the heat conduction  into the interior, or both—as a consequence of the dynamic evolution of the surface and bulk oxide  properties. The condensed phases for the oxide scale of the ZrB2-SiC composites are ZrO2 and SiO2  under conditions 1 and 2  respectively. During  the  initial heating process  in aerothermal  test  environments, SiC underwent passive oxidation, resulting in the formation of SiO2 that was mostly  distributed on the surface, since the temperature has not reached the passive-to-active oxidation  transition. Following that, active oxidation of SiC would occur with the prolongation of the heating  time, accompanied by a further increase of the surface temperature up to a steady state. It should be  noted that the steady state temperature is much higher than that of the passive-to-active oxidation  transition for SiC in the case of present tests [19]. Obviously, the active oxidation of SiC has minor  effect on  the change of  the surface composition when  the specimen achieved  the steady state.  Therefore, the temperature jump is not directly associated with a passive-to-active transition of SiC,  since the active oxidation of SiC for ZS30 were also observed under both conditions, as shown in  Figure 7, where no large temperature jump occurred. Note that silica volatilization becomes significant  as temperature increases up to a steady state temperature (~1900 °C). Consequently, the amount of  silica in the surface layer decreases gradually with the prolongation of the testing time which would  lead to an increase of catalytic efficiencies for oxygen recombination, because silica surfaces are  known to have relatively low catalytic efficiencies for oxygen recombination compared to zirconia  surfaces [28,29]. Therefore, the sample will experience a higher chemical heating contribution to the  heat flux for a given free stream condition, especially under high enthalpy conditions. A pronounced   increment in surface temperature within a short time interval will be observed when the surface silica has   reduced   to a certain   level. The   thickness of oxide scale   increased with   increasing   time before   the   temperature jump, and hence the heat conduction into the interior decreases because of lower heat  conductivity of oxide scale, compared to that of the base material. In addition, heat conductivity of the  oxide scale was also influenced by microstructural evolution during the oxidation process. These will  lead to a rise in surface temperature. However, the oxide scale is adherent to the base material, and  passive oxidation of UHTCs was detected before a temperature jump in the present case, and its heat       \\x0c', 'Materials 2013, 6   1741   conductivity changed gradually, which would not have caused a large temperature jump in a short   time. To verify the strong influence of the composition on the surface temperature, ablation behavior of the   sample containing AlN was carried out as a comparison. The surface temperature of this composite is  significantly higher (by ~700 °C)   the oxide of AlN (alumina)   is a    than   those without AlN since   well-known high catalycity material [27,30]. Therefore, the different chemical heating of the surface  oxide  layer contributes mostly  to  the  large  temperature  jump, which depends on  the material  composition and testing conditions.  In the present study, the temperature jump always occurred at the edge of the sample and then  propagated to other areas, especially around the sample periphery (as shown in Figure 2), because the  heat flux and shear force at the edge are higher than those in the center for the cylinder shape model.  This might cause the fast loss of silica around the sample periphery, leading to a significant increase in  surface  temperature as a result of substantial  increases  in  the chemical heating component  that  delivered to the surface.   Generally, the temperature jump occurred at high temperatures, above 1900 °C, under the gas  pressure used in this study. The exact value for temperature was hard to measure during the testing  because of experimental constraints. The related test results indicate that the variation of pressure in a  certain range has a minor effect on the transition of temperature jump. Additionally, the time for the  jump was dependent on  the  initial steady state  temperature, and decreased with  the  increasing  temperature, since the change of the surface composition at higher temperature became faster, so  played an important role in the chemical heating contribution. The temperature jump always occurred  at high temperature (e.g., >1900 °C), in several minutes, and the time to trigger the spontaneous  temperature jump was shortened as the surface temperature increased because the loss of silica  increased rapidly with increasing temperature. In contrast, the temperature jump did not occur at lower  temperatures, since the oxide scale was always covered by a stable silica glass layer during the entire  testing period. Meanwhile, small temperature fluctuation and good reproducibility were observed on a  series of sample tests at low heat flux condition, i.e., with a lower surface temperature. Furthermore,  the time to initiate the spontaneous temperature jump was affected by the thickness of the silica glass.  The surface scale containing silica with a higher thickness needed more time to consume. The steady  state time during the testing was correlated to the loss rate of SiO2 and the amount of pre-existing SiO2  at the beginning of steady state. Lower SiC content produced less SiO2 during the initial heating  process, leading to the temperature jump and a higher surface temperature in the current tests. The  different steady state surface temperatures achieved for the different compositions were related to the  steady state energy balance at the surface, as determined by radiation, catalysis, oxidation, and  convection. The temperature jump is a transient phenomenon related to a transition in Si species  related surface chemistry that changes the surface heating rate under constant free stream conditions.  The molten oxide was found after the temperature jump (shown in Figure 2) and then quickly spread  to the entire sample surface. The temperature increase of the surface layer also induced the temperature  increase of the interior reaction region. This in turn led to formation of high pressure gas phases (e.g.,  B2O3, CO, and SiO) within the oxide layer, causing the rupture of the oxide scale. The ejection of molten  oxide was detected in video images. A large number of pores were generated after the ejection of molten  oxide, enabling the free release of the formed gas phase products, and then sample surface reached a new  steady state. Poor adherence and decreased thermal contact of exterior scale to the underlying material       \\x0c', 'Materials 2013, 6   1742   may contribute to large temperature gradients through the oxide scale after the temperature jump. The  outside scale acted as a thermal barrier coating, and this caused the decrease of the oxidation reaction rate  of the inner material. Therefore, the outside scale provides an efficacious protective ablation barrier,  although it is not very effective in limiting the inward diffusion of oxygen into the inner bulk.  The performance of ZrB2-based UHTC composites was significantly affected by  the surface  temperature, which seems associated with surface and bulk oxide layer properties that strongly depend  on the material composition and testing conditions. The microstructural evolution of the oxide scale  and the change of surface chemistry as the testing time proceeded caused the temperature to rise  sharply as a result of the thermal response change under high heat flux condition, especially with high  enthalpy. The ablation and oxidation properties of UHTCs are directly correlated with the given  service environment, as well as the material itself. The furnace oxidation testing method is cheaper  compared with other methods, such as arc-jet testing and high frequency plasma testing, and allows the  sensitive control of temperature, oxidation time, oxidizing atmosphere, and pressure, whereas it cannot  reflect the thermal response of the service environment on the material. It is worth noting that the  oxidation performance of UHTCs is strongly dependent on the specific oxidizing environment. For  example, the addition of Ta-containing compounds is beneficial to the oxidation resistance of UHTCs  at temperatures below 1700 °C, whereas it has a detrimental effect on the oxidation resistance at  temperatures of 1800 °C or above [31]. Oxidation testing results involving both furnace oxidation  testing and plasma wind tunnel testing (or arc-jet testing) indicated that the oxidation resistance of  UHTCs is strongly dependent on the oxidation temperature [8-16]. The present results showed that the  oxidation performance of UHTCs can be  improved by reducing  the surface  temperature of  the  material, based on the composition optimization in the given oxidizing environment. However, the  effect of material composition on thermal response under a given aerodynamic heating condition is  unclear and additional testing will be required in this respect in the future. Therefore, it is advisable to  balance  the  thermal response and oxidation resistance of  the UHTCs against  temperature when  choosing materials for extreme oxidizing environment.    5. Summary   ZS10 displayed significant ablation after tests under the two conditions, whereas ZS30 exhibited an  excellent ablation resistance and configurational stability. The significant difference  in ablation  behavior was mostly attributed to different surface temperature. ZS10 showed the highest surface  temperature (>2300 °C) while ZS30 exhibited the lowest surface temperature (<2000 °C) in both  conditions. The marked difference in surface temperature was presumably caused by the different  chemical component of heat flux delivered to the surface, which strongly depended on the material  composition and aerothermal heating environment. This in turn affected its oxidation performance. It is  advisable to balance the thermal response and oxidation resistance of the UHTCs against temperature  when choosing materials for extreme oxidizing environment.   Acknowledgments   This work was supported by the National Science Foundation of China (51202048, 51072042   and 91216301).       \\x0c', 'Materials 2013, 6   References   1743   1.   2.   3.   temperature   structural   for ultra high   Upadhya, K.; Yang,  J.M.; Hoffman, W.P. Materials  applications. Am. Ceram. Bull. 1997, 72, 51-56.  Opeka, M.M.; Talmy, I.G.; Zaykoski, J.A. Oxidation-based materials selection for 2000 °C+  hypersonic aerosurfaces: Theoretical considerations and historical experiences. J. Mater. Sci. 2004,  39, 5887-5904.  Opeka, M.M.; Talmy, I.G.; Wuchina, E.J.; Zaykoski, J.A.; Causey, S.J. Mechanical, thermal, and  oxidation properties of hafnium and zirconium compounds. J. Eur. Ceram. Soc. 1999, 19, 2405-2414.  4. Monteverde, F. 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Recombination of atomic oxygen on sintered zirconia  at high temperature in non-equilibrium air plasma. Mater. Chem. Phys. 2010, 123, 40-46.  30. Li, G.; Han, W.B.; Zhang, X.H.; Han, J.C.; Meng, S.H. Ablation resistance of ZrB2-SiC-AlN  ceramic composites. J. Alloys Compd. 2009, 479, 299-302.  31. Hu, P.; Zhang, X.H.; Han, J.C.; Luo, X.G.; Du, S.Y. Effect of various additives on the oxidation  behavior of ZrB2-based ultra high temperature ceramics at 1800 °C. J. Am. Ceram. Soc. 2010, 93,  345-349.   © 2013 by the authors; licensee MDPI, Basel, Switzerland. This article is an open access article  distributed  under  the  terms  and  conditions  of  the Creative Commons Attribution  license  (http://creativecommons.org/licenses/by/3.0/).       \\x0c', \"Copyright of Materials (1996-1944) is the property of MDPI Publishing and its content may not be copied or  emailed to multiple sites or posted to a listserv without the copyright holder's express written permission.  However, users may print, download, or email articles for individual use.  \\x0c\"]"
},{
  "_id": 53,
  "PDF": "Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride.pdf",
  "Text": "['http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  Fei Peng, Yolande Berta, and Robert F. Speyera)  School of Materials Science and Engineering, Georgia Institute of Technology,  Atlanta, Georgia 30332-0245  (Received 4 November 2008; accepted 9 February 2009)  The oxidation resistances of ZrB2 containing SiC, TaB2, and TaSi2 additions of various  concentrations were studied using isothermal thermogravimetry at 1200, 1400, and  1500  \\x0e  C, and specimens were further characterized using x-ray diffraction and electron  microscopy. Increasing SiC concentration resulted in thinner glassy surface layers as well  as thinner ZrO2-rich underlayers deficient in silica. This silica deficiency was argued to  occur by a wicking process of interior-formed borosilicate liquid to the initially-formed  borosilicate liquid at the surface. Small (3.32 mol%) concentrations of TaB2 additions  were more effective at increasing oxidation resistance than equal additions of TaSi2.  The benefit of these additives was related to the formation of a zirconium-tantalum boride  solid solution during sintering, which during oxidation, fragmented into fine particles of  ZrO2 and TaC. These particles resisted wicking of their liquid/glassy borosilicate  encapsulation, which increased overall oxidation resistance. With increasing TaB2 or  TaSi2 concentration, oxidation resistance degraded, most egregiously with TaB2  additions. In these cases, zirconia dendrites appeared to grow through the glassy layers,  providing conduits for oxygen migration.  I.  INTRODUCTION  Transition metal borides  including ZrB2, TaB2,  and  HfB2 are of interest for their ultra-high melting temperature (>3000 C), high hardness and strength, and high thermal and electrical conductivities.1-3 They are candi \\x0e  dates for leading edges on reentry vehicles; their survival  against atmospheric frictional heating is dependent on a  combination of refractory ability and the ability to dissi pate  heat  through  thermal  conduction  away  from the  leading edge and radiant emission to the cold ambient.  Engineering of  these ceramics  for oxidation resistance  has focused on a two-phase microstructure of ZrB2 and  SiC,  in which a borosilicate viscous liquid with interdis persed ZrO2 forms as a passivating surface layer.  Oxidation of single-phase ZrB2 does not  form a pro tective surface layer since B2O3 is volatile (boiling point,  i.e., 1 atm vapor pressure, of B2O3 is 1860  \\x0e  C). Oxidation  heat  treatments of ZrB2 + 20 vol% SiC at 1200  \\x0e  C and  below have shown weight gain no less extensive than  those of specimens composed of ZrB2 alone. However, C, a borosilicate coating forms.4,5 Given the  above 1200  \\x0e  high volatility of boron oxide,  the borosilicate glass sur face  coating might be  expected to become more of  a  near-pure fused silica coating with increasing tempera ture. One investigation has shown that the boron content  of the oxide layer after heating to 1500 C for 30 min is less than 1 wt%.6 However, B2O3 vapor pressure is suppressed by its entering into solution with SiO2. Further,  \\x0e  standard  glass-forming  practice melts,  homogenizes,  and fines borosilicate (e.g., Pyrex) glass compositions at  1550-1600 C with residence times of several hours to days.7 However, far less of the boron oxide component of  \\x0e  this  liquid would be exposed to the liquid-vapor  inter face, where it could volatilize, as would occur  in a thin  liquid/amorphous coating.  Opila et al. showed that a ZrB2-20 vol% SiC compo sition exposed to ten 10 min oxidation cycles at 1327 C developed protective oxide scales: 30 mm at and 1627 C and 150 mm at 1627 C.8 Thermal cycling at  \\x0e  1327  \\x0e \\x0e  \\x0e  1927  C resulted in an oxide  layer  thickness of over  1 mm. The 1627  \\x0e  C surface oxide coating was identified  (via energy dispersive spectrometry) to be silica. Under neath this  coating was  a  region of ZrO2 dispersed in  silica, which in turn was above a region of ZrB2 depleted  of SiC. This region was argued to have resulted from active oxidation of SiC to form SiO(g).8,9 Opeka et al. have suggested that formation of SiO(g) could build up to  pressures exceeding ambient,  facilitating rupture of  the  protective glass layer, resulting in a cyclic protective/ non-protective scale-forming sequence.10  In  low oxygen  partial  pressures,  formed  by CO(g)/  CO2(g) mixtures, ZrB2 oxidized to ZrO2(s) and a volatile  a)Address all correspondence to this author.  e-mail: robert.speyer@mse.gatech.edu  DOI: 10.1557/JMR.2009.0216  J. Mater. Res., Vol. 24, No. 5, May 2009  © 2009 Materials Research Society  1855  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  boron oxide, and SiC oxidized to carbon monoxide and scale.6  SiO(g). This  left a non-protective (porous) ZrO2  Han et al. found that for ZrB2-20 vol% SiC, the thickness  of the oxide layer increased and oxidation resistance decreased with decreasing oxygen partial pressure.11 Karls dottir  et  al.  showed that  zirconia  is deposited on the  surface of  the glassy coating during oxidation at 1550  and 1700  \\x0e  C by convection of a boron-zirconium-silicon  oxide liquid phase which evolves volatile B2O3 at surface, precipitating ZrO2.12 Talmy et al. investigated additions of Cr-, Ti-, Nb-, V-,  the  and Ta-borides to ZrB2-25 vol% SiC, and found that all  additions (all of which formed solid solutions with ZrB2  after  sintering)  improved oxidation resistance over  the  base composition, with TaB2 additions being the most effective.13 It was found that  improved oxidation resis tance correlated with increasing cation field strength (defined as Z/r2, where Z is the valance of the cation and r is  the ionic radius) of the added diborides. In a borosilicate  or  silicate glass with transition metal cations,  the ten dency toward liquid immiscibility is known to increase  with  increasing  cation  field  strength  of  the  transition  metal. This phase separation has been argued to result in increased viscosity,14 which has been correlated to reduced oxygen diffusion rates.10  Opila et al. found that TaSi2 additions in the form of a  ZrB2-20 vol% SiC-20 vol% TaSi2 composition showed a  lower oxidation rate after cyclic oxidation at 1627 C than a ZrB2-20 vol% SiC composition.15 Improved oxidation resistance was related to evidence of phase sepa \\x0e  ration in the amorphous surface layer. The composition  containing TaSi2  showed  rapid  consumption  as  com pared  to ZrB2-20  vol% SiC compositions  exposed to  similar oxidation heat  treatments at 1927  \\x0e  C. This was  attributed to melting of Ta2O5  (1785  \\x0e  C) and/or com pounds of Ta2O5 and ZrO2. Talmy et al. demonstrated  enhanced oxidation resistance from adding 8-30 vol% Ta5Si3 to ZrB2.16 A purported general advantage of tantalum compound additions is that tantalum can stabilize  zirconium oxide,  circumventing  the  tetragonal/mono clinic phase transformation, whose volume change can create fissures in the oxide scale.17  Zhang et al. have reported that a ZrB2-SiC composi tion with 10 vol% LaB6 additions showed good oxidation  resistance at 2400  \\x0e  C (oxyacetylene torch);  lanthanum  stabilizes the tetragonal form of zirconia, and the La2Zr2O7  along with zirconia forms a coherent and compact oxide surface scale.19 Fahrenholtz et al. have recently shown the benefits of tungsten additions to ZrB2.18 The tungsten functions as a sintering aid to the surface zirconia which forms  via oxidation, decreasing the oxygen permeability through  this surface layer. In our previous work,20  the oxidation resistances of  ZrB2  containing SiC, TaB2,  and TaSi2  additives were  studied  using  scanning  thermogravimetry  (3  \\x0e  C/min)  over  the range 1100-1550  \\x0e  C.  It was  shown that SiC  additions to ZrB2 improved oxidation resistance, as did  TaB2, and to a greater extent, TaSi2 additions.  In this  investigation, relative oxidation resistances of these com positions and their causes were studied in more detail  via isothermal  thermogravimetry studies at 1200, 1400,  and 1500  \\x0e  C.  II. EXPERIMENTAL PROCEDURE  Commercially-available powders were used for  raw  materials. The major  crystalline  phase(s),  grade,  and  suppliers  are  listed  for  each  powder  in Table  I. The  particle sizes of commercially-available TaB2 and TaSi2  were deemed too large for pressureless sintering. Hence,  sedimentation-based selection was used to obtain finer  particles: Powders were dispersed in ethanol using an  ultrasonicator  (FS-14 Solid State Ultrasonicator, Fisher  Laboratory  Equipment Division,  Pittsburgh,  PA)  for  10 min. The mixture was allowed to settle in ethanol for  1 h. The top 7 cm of  fluid was  then extracted using a  pipette. Based on laser particle size analysis (Model LS  13 320, Beckman Coulter, Fullerton, CA), decanted particles had a d50 of 1.1 mm for TaB2 and 1.7 mm for TaSi2. The decanted suspensions were dried in a beaker on a  hot plate.  The compositions of synthesized powder mixtures are  given in Table II. The powder mixtures were suspended  in methanol, and mixed in a ball mill for 24 h, using B4C  as media. The milled powders were then dried in static  air at 75  \\x0e  C. The powder mixtures were then ball milled  again in water with dissolved polyvinyl alcohol  (PVA,  Celanese Ltd., Dallas, TX), polyethylene glycol  (PEG,  Alfa Aesar, Ward Hill, MA), and Darvan 821A (R.T.  Vanderbilt Company Inc., Norwalk, CT), using B4C as  media for 8 h. PVA functioned as a binder with PEG  functioning as a plasticizer, and Darvan 821A served as  a dispersing agent. The highly viscous suspension after  this milling step was dried in an oven at 75  \\x0e  C, and then  sieved using a 60 mesh screen.  Approximately  400 mg  of  powder  were  uniaxi ally pressed into cylindrical pellets using a pressure of  TABLE I. Raw material characteristics.  Phases  Particle size  Supplier  ZrB2  ZrB2  d50 = 2.20 mm  Grade B, H. C.  Starck, GmbH  B4C  stoichiometric  B4C a-SiC  d50 = 0.8 mm  Grade HS, H. C.  Starck, GmbH  SiC  d50 = 0.88 mm  Grade 8S490NDP,  Superior Graphite,  Chicago, IL  TaB2  TaB2, Ta3B4  <43 mm  ESPI Metals, Ashland,  OR  TaSi2  TaSi2  <43 mm  Cerac Inc., Milwaukee, WI  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1856  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  117 MPa, holding for 1 min. The pellets were loaded into  latex encapsulants which were in turn evacuated. These  were  cold isostatically pressed (CIP)  at 345 MPa  for  1 min. This was followed by a binder-removal heat treat ment of 0.25  \\x0e  C/min to 500  \\x0e  C under vacuum (\\x184 Pa).  Fifteen pellets were fabricated for each composition.  These pellets were  fired in a graphite  tube  furnace  (Model M11, Centorr Vacuum Industries  Inc., Nashua,  NH) under  flowing argon, using graphite to \\x184  setters. The  furnace was  initially  evacuated  Pa  (roughing  pump)  and backfilled with argon. The  typical heating  schedule was 50  \\x0e  C/min to 2000  \\x0e  C,  soaking for 1 h,  and then cooling at 40  \\x0e  C/min to room temperature. The  pellets were then hot  isostatically pressed (HIP, Ameri can Isostatic Press, Columbus, OH) at 1800  \\x0e  C for 30  min under an argon gas pressure of 207 MPa. The den sities of unfired pellets were determined from measured  dimensions and mass;  the densities of pressureless sin tered  and  post-HIPed  pellets were  determined  using  Archimedes’ method. All  specimens were 100% dense  based on theoretical densities calculated from the rule of  mixtures.  All of  the  surfaces of  all post-HIPed samples were  ground away using 320 grit SiC grinding paper (Buehler,  Lake Bluff, IL), and the resulting pellet dimensions were  measured with calipers. The oxidation behaviors were  then investigated using thermogravimetric analysis (TG,  Model STA 409, Netzsch, Exton, PA, with an Innovative  Thermal Systems  interface, Atlanta, GA). The samples  were supported on alumina chips which filled an alumi na  crucible,  to minimize  the  contact  between  sample  TABLE II. Sample compositions in mol%.  Code  ZrB2  B4C  SiC  TaB2  TaSi2  ZBS2  77.39  7.27  15.34  0  0  ZBS6  70.20  6.60  23.20  0  0  ZBS10  64.25  6.04  29.70  0  0  ZBS14  57.53  5.40  37.06  0  0  ZBS18  50.78  4.78  44.45  0  0  ZBS22  43.80  4.12  52.08  0  0  ZBS26  37.55  3.53  58.92  0  0  ZBSTB1  63.14  5.62  27.91  3.32  0  ZBSTB2  59.82  5.62  27.91  6.65  0  ZBSTB3  56.49  5.62  27.91  9.97  0  ZBSTB4  53.17  5.62  27.91  13.29  0  ZBSTB5  49.85  5.62  27.91  16.61  0  ZBSTS1  63.14  5.62  27.91  0  3.32  ZBSTS2  58.82  5.62  27.91  0  6.65  ZBSTS3  56.49  5.62  27.91  0  9.97  ZBSTS4  53.17  5.62  27.91  0  13.29  ZTBS1-1  59.82  5.62  27.91  6.65  0  ZTBS1-5  59.82  5.62  27.91  3.32  3.32  ZTBS1-9  58.82  5.62  27.91  0  6.65  ZTBS2-1  53.17  5.62  27.91  13.29  0  ZTBS2-5  53.17  5.62  27.91  6.65  6.65  ZTBS2-9  53.17  5.62  27.91  0  13.29  TABLE III. Phases identified via XRD of as-synthesized and oxidized specimens at indicated temperatures (“ss” indicates solid solution, “tr” indicates trace quantities, M and O indicate the monoclinic and orthorhombic forms of zirconia, respectively).a  Code  As-fabricated  1200  \\x0e  C  1400  \\x0e  C  1500  \\x0e  C  ZBS2  ZrB2  M-ZrO2  M-ZrO2  M-ZrO2  ZBS6  ZrB2  M,O-ZrO2  M-ZrO2  M,O-ZrO2  ZBS10  ZrB2, SiC(tr)  M-ZrO2  M-ZrO2, SiC  M,O-ZrO2  ZBS14  ZrB2, SiC  M-ZrO2, SiC  M-ZrO2, SiC  M-ZrO2  ZBS18  ZrB2, SiC  M-ZrO2, SiC, ZrB2  M-ZrO2, SiC, ZrB2  M-ZrO2, SiC, ZrB2  ZBS22  ZrB2, SiC  M-ZrO2, SiC, ZrB2  M-ZrO2, SiC, ZrB2  M-ZrO2, SiC, ZrB2  ZBS26  ZrB2, SiC  M-ZrO2, SiC  M-ZrO2, SiC  M-ZrO2, SiC, ZrB2  ZBSTB1  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC  M-ZrO2, TaC  M,O-ZrO2, TaC  ZBSTB2  ZrB2-TaB2(ss) SiC  M,O-ZrO2, TaC  M-ZrO2, TaC  M,O-ZrO2, TaC  ZBSTB3  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC  M,O-ZrO2, TaC  M,O-ZrO2, TaC  ZBSTB4  ZrB2-TaB2(ss), SiC  M-ZrO2, TaC, SiC, ZrC(tr)  M,O-ZrO2, TaC, TaB2  M,O-ZrO2, TaC, TaB2  ZBSTB5  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC, TaZr2.75O8  M,O-ZrO2, TaC, TaB2  M,O-ZrO2, TaC  ZBSTS1  ZrB2-TaB2(ss), SiC  M-ZrO2, TaC, ZrC(tr)  M,O-ZrO2, TaC  M,O-ZrO2 TaC  ZBSTS2  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC  M,O-ZrO2, TaC, SiC,  ZrB2-TaB2(ss)  M,O-ZrO2, TaC, ZrB2-TaB2(ss)  ZBSTS3  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC ZrB2-TaB2(ss), ZrC(tr)  M-ZrO2, TaC, ZrB2-TaB2(ss),  ZrC(tr)  M,O-ZrO2, TaC ZrB2-TaB2(ss)  ZBSTS4  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC, ZrB2-TaB2(ss), ZrC(tr)  M-ZrO2, TaC, TaB2, ZrC(tr)  M,O-ZrO2, TaC, TaB2  ZTBS1-1  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC  M-ZrO2, TaC  M,O-ZrO2, TaC  ZTBS1-5  ZrB2-TaB2(ss), SiC, TaSi2(tr)  M-ZrO2, TaC, TaB2  M-ZrO2, TaC, SiC,  ZrB2-TaB2(ss)  M-ZrO2, TaC, ZrB2-TaB2(ss)  ZTBS1-9  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC  M,O-ZrO2, TaC, SiC,  ZrB2-TaB2(ss)  M,O-ZrO2, TaC  ZTBS2-1  ZrB2-TaB2(ss), SiC  M-ZrO2, TaC, SiC, ZrC(tr)  M,O-ZrO2, TaC, TaB2  M,O-ZrO2, TaB2, TaC  ZTBS2-5  ZrB2-TaB2(ss), SiC  M-ZrO2, TaC, TaB2  M-ZrO2, TaC, TaB2, ZrC(tr)  M,O-ZrO2, TaC, TaB2  ZTBS2-9  ZrB2-TaB2(ss), SiC  M,O-ZrO2, TaC, ZrB2-TaB2(ss), ZrC(tr)  M-ZrO2, TaC, TaB2, ZrC(tr)  M,O-ZrO2, TaC, TaB2  a Some compositions are indicated by two different codes. This was done to display compositional trends in logical groupings.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1857  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  surfaces and alumina. Samples were exposed to flowing  air from a compressed air tank with a flow rate of 0.1 L/  min. The  flow rate was maintained  via  a mass  flow  controller  (Model GFC 17, Aalborg, Orangeburg, NY).  Specimens were heated to soak temperatures of 1200,  1400, or 1500  \\x0e  C and held for 4 h. For the 1200  \\x0e  C soak,  specimens were heated at 50  \\x0e  C/min to 950  \\x0e  C, 30  \\x0e  C/  min to 1100  \\x0e  C, 10  \\x0e  C/min to 1170  \\x0e \\x0e  C, and 5  \\x0e  C/min to  the soak temperature. For  the 1400  C soak, specimens  were heated at 50  \\x0e  C/min to 1150  \\x0e  C, 30  \\x0e  C/min to  1300  \\x0e  C, 10  \\x0e  C/min to 1370  \\x0e \\x0e  C, and 5  \\x0e  C/min to the  soak temperature. For the 1500  C soak, specimens were  heated at 50  \\x0e  C/min to 1250  \\x0e  C, 30  \\x0e  C/min to 1400  \\x0e  C,  10  \\x0e  C/min to 1470  \\x0e  C, and 5  \\x0e  C/min to 1500  \\x0e  C. Dis played data is truncated to the start of the soak tempera ture. To evaluate repeatability, 3-4 TG oxidation heat  treatments were performed on each composition. Dis played traces are those considered most representative.  Crystalline phases in the samples were identified us ing X-ray diffraction (XRD, Model X’Pert PRO Alpha 1, PANalytical, The Netherlands). Scans were recorded at room temperature over a 2y range of 10-80  \\x0e  at a scan  rate of 0.01 deg/s. Prior to oxidation heat treatment, spec imen surfaces were ground to expose specimen interior  regions for XRD analysis. The XRD of oxidized speci mens were  taken from unaltered surfaces. The micro structures of oxidized samples were investigated using  scanning electron microscopy (SEM, LEO 1530, Carl  Zeiss SMT,  Inc., Thornwood, NY) and energy disper sive spectrometry (EDS, Oxford Pentafet detector with  FIG. 1. Microstructures of fracture surfaces of two ZBS compositions. EDS data and compositional contrast  imply dark round regions are B4C,  slightly lighter-shaded, sometimes elongated, grains are SiC, and light-shaded regions are ZrB2.  FIG. 2. TG of ZrB2-B4C specimens with varying amounts of SiC (ZBS series), soaked at  three different  temperatures. Numbers in the plots  indicate SiC content in mol%. Data sets consist of approximately 2000 data points; the symbols on the traces are for curve identification purposes.  This is the case for all TG traces in this paper.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1858  \\x0c', 'F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  ultrathin window, Oxford Instruments, Oxfordshire, UK).  mens evaluated by XRD, SEM, and TEM were surface Specimen cross sections were formed via impact-formed  oxidized in the TG as previously described.  fracture surfaces, with the oxidized surfaces mounted on  the SEM stub to be parallel with the beam axis. Speci mens were coated with gold (sputtering for 2 min) to form  a conductive surface.  III. RESULTS  Phases  identified in as-fabricated specimen interiors,  Thin sections of specimen cross sections were prepared  as well as phases identified from surfaces after oxidation  by using a focused-ion-beam milling system (Nova Nano heat  treatments, are listed in Table III. For  the as-fabri lab 200 FIB/SEM system, FEI Corp., Hillsboro, OR). Approximately 2 mm of Pt was deposited on the top of  the samples to avoid ion beam damage. The samples were  milled through the specimen cross section using a gallium  ion beam at 30 kV, 30 pA to 20 nA, until reached \\x18100 nm. These  thicknesses  specimen  specimens were  cated  specimens,  the  boron  carbide  additive was  not  detected. TaB2 formed a solid solution with ZrB2 as evidenced by distinct shifts in the 2y values of the ZrB2 peaks. This also occurred with TaSi2 additions.  Figure 1 shows  the microstructures of  fracture  sur faces of densified ZBS samples of  two different SiC  analyzed for microstructure on a JEOL 4000EX 400kV  concentrations. The darker phase is SiC, and the lighter  high resolution transmission electron microscope (TEM,  JEOL USA, Peabody, MA) with a point-to-point resolu is ZrB2, as would be expected based on atomic weights,  and  as  implied  by  the  EDS  spectra  of  the marked  tion of 0.18 nm, and a Hitachi HF-2000 200 kV field  regions. SiC appeared as a continuous phase in the spec emission gun TEM (Hitachi High Technologies America,  imen with 52.1 mol% SiC, and as an isolated phase in  Inc., Pleasanton, CA) with EDS capabilities (Noran de the specimen with 15.3 mol% SiC. The occasional round  tector with ultrathin window,  for detection of elements  darkest regions (e.g., region B in the upper micrograph)  down to boron, Thermo Scientific, Madison, WI). Speci are interpreted to be B4C, based on comparisons of EDS  FIG. 3. Cross  section microstructures of ZBS series  samples with varying concentrations of SiC, oxidized in flowing air  at  the  indicated  temperatures. Energy dispersive spectroscopy spectra corresponds to the marked regions in the microstructures. EDS peaks identified as gold  correspond to conductive coatings applied to specimen surfaces.  http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  J. Mater. Res., Vol. 24, No. 5, May 2009  1859  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  spectra and the dark shading consistent with light ele ments.  In general, boron has not been reliably detected  using the available EDS detector.  After  oxidation  heat  treatments  of  the ZBS series,  zirconia in its  stable monoclinic form was detected by  XRD (orthorhombic zirconia was detected in trace quan tities  at  1500  \\x0e  C). SiC and ZrB2 were  also  detected  for some of  the higher SiC-containing samples. For  the  various  series containing TaB2 and/or TaSi2, TaC was  detected (on occasion,  trace quantities of ZrC were also  detected). TaB2 additions resulted in no apparent shifting in ZrO2 2y peak positions, nor were increases in peak intensity of orthorhombic-ZrO2.  there noticeable  Figure 2 shows TG traces of  the ZBS series at  three  different soaking temperatures. Mass increases after  the  4 h soaks were less extensive with increasing SiC con tent. The rates of weight gain were decreasing with time,  with the  exception of  low concentrations of SiC (i.e.,  <27.9 mol% SiC) at 1200  \\x0e  C.  Figure 3 shows  the cross  section microstructures of  ZBS  specimens with  varying  SiC concentrations,  in  which a glassy surface coating is seen covering an oxide  underlayer. Based on comparisons of EDS spectra, mi crostructure morphology,  and XRD results,  the  oxide  underlayer  is  interpreted  to  be  zirconia  particles  sur rounded  by  porosity  and  some  remnant  borosilicate  glass. For  the 1500  \\x0e  C heat  treatment, both the glassy  and  oxide  underlayer  thicknesses  decreased with  in creasing SiC concentration. EDS patterns show zirconi um along with silicon in the glassy surface layer for the  specimen with 15.3 mol% SiC, but only silicon in the  glassy layer for the specimen with 29.7 mol% SiC.  The  glassy  surface  layer  and  the  oxide  underlayer  thicknesses as a function of  soak temperature and SiC  concentration, based on measurements on SEM micro graphs, are depicted in Fig. 4(a). Figure 4(b) summarizes  the weight gains (based on weight  losses after 4 h soaks  determined from Fig. 2) as a function of SiC concentra tion and soak temperatures.  Figure 5 shows TG traces of  the ZBSTB series (with  varying TaB2 concentrations). For specimens soaked at  1200  \\x0e  C, higher TaB2 concentrations resulted in increased  oxidation resistance, while for soaks at 1500  \\x0e  C, increas ing  TaB2  concentrations  showed  diminishing  oxida tion resistance. The weight change was low (lower than  many compositions  soaked at 1200  \\x0e  C), and relatively  unchanged with composition,  for  specimens  soaked at  FIG. 5.  Isothermal TG traces of ZrB2-B4C-SiC specimens containing  varying amounts of TaB2 (ZBSTB series). Solid curves correspond to  1200  \\x0e  C heat  treatments, dashed curves correspond to 1400  \\x0e  C heat  treatments, and dot-dashed curves correspond to 1500  \\x0e  C heat  treat ments. Numbers in the figure indicate mol% of TaB2.  FIG. 4. Effects of oxidation heat treatments (in the TG) for ZBS compositions of varying SiC content. (a) Layer thicknesses estimated from SEM  micrographs of specimen cross sections. Solid lines indicate the thicknesses of the oxide underlayers. Dashed lines indicate the thicknesses of the  glassy layers adjacent  to the surfaces. Temperatures marked in the figure refer  to the isothermal  soak temperatures.  (b) Mass changes after  soaking at  the indicated temperatures for 4 h in the TG.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1860  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  1400  \\x0e  C. Higher  (13.29  and  16.61 mol%)  concentra tions of TaB2 resulted in marked points of acceleration C (at \\x1860 in mass min for 16.61 mol% TaB2, and at \\x18120 min for 13.29 increase during oxidation at 1500 mol% TaB2).  \\x0e  Figure 6 shows glassy surface layers and oxide under layers covering unreacted nonoxide phases for ZBSTB5.  EDS spectra show Si, Zr, and Ta in the oxide underlayer.  No tantalum was confirmed by EDS in the glassy layer.  Figure 7 shows  a TEM micrograph of  the amorphous  surface layer with a rectangularly-shaped zirconia crys tal;  such crystals were observed occasionally in TEM  analysis of the glassy surface layer. No tantalum is indi cated in the glass,  though major Ta peaks overlap with  Si and Cu, which may mask small concentrations. There  was no evidence of phase  separation in the glass. As  shown in Fig. 8, the oxide underlayer consists of zirconia  and TaC (dark round) particles which are incompletely coated by a glassy phase. The TaC crystals were \\x1850  nm—substantially smaller  than the ZrO2  crystals. The  glassy phase  shows  faint  evidence of  tantalum, while  the  zirconia  crystals  show slightly  better  evidence  of  tantalum. Figure 9(a) displays glassy surface layer and  oxide underlayer  thicknesses, while Fig. 9(b)  summa rizes weight changes for the ZBSTB series.  Figure 10 shows TG traces of  the ZBSTS specimens  (with varying TaSi2  content). Figure 11(a)  shows  the  glassy and oxide  layer  thicknesses  for  these  composi tions, and Fig. 11(b)  summarizes  the oxidation weight  gain measurements. Oxide  underlayer  thicknesses  de creased with initial  increases  in TaSi2 content. Glassy  layer  thicknesses were relatively constant at 1400 and  1500  \\x0e  C; at 1200  \\x0e  C,  the layer did not  form for TaSi2  contents less than 6.65 mol%. As with TaB2 additions,  the effects of TaSi2 on oxidation resistance reversed at  1500  \\x0e  C; TaSi2 additions above 3.32 mol% resulted in a  decreased oxidation resistance,  though not nearly as ex tensive  as  observed with  increasing TaB2. Figure  12  shows the effect of substitution of TaSi2 for TaB2 (ZTBS  series). For  these compositions, TaSi2 is shown to be a  more  effective  additive  for  oxidation  resistance  than  equimolar additions of TaB2.  The  appearances  of  the  oxide  underlayers  in  the  ZBS10 (no tantalum additive), ZBSTB5 (TaB2 added),  FIG. 6. Cross section microstructures and selected EDS spectra of ZBSTB5 for two soak temperatures. Although tantalum was clearly detected  by EDS in the oxide underlayers, its complete absence in the glassy layer cannot be confirmed because of EDS peak overlap with other elements.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1861  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  and ZBSTS4 (TaSi2 added) series are compared in Fig. 13  for the 1400  \\x0e  C heat treatment, and Fig. 14 for the 1500  \\x0e  C  heat treatment. Zirconia particles were substantially smal ler in the specimens containing TaB2 or TaSi2 additions.  A relatively substantial amount of glassy phase was seen  interdispersed in the oxide underlayer in samples contain ing TaSi2 as a batch additive.  IV. DISCUSSION  In the ZBS series, ZrB2 and SiC form distinct phases  (Table III and Fig. 1). While B4C was not detected by  XRD, its presence as a distinct phase is implied by SEM/  EDS. TG traces (Fig. 2) of  these compositions at 1400  and 1500  \\x0e  C show a decreasing weight  change  slope  with time,  indicative of  the buildup of a barrier coating  against  the diffusion of  reactant and product gases. At  1200 tions (\\x1415.3 mol% SiC, Fig. 2), C, this was not observed for  \\x0e  low SiC concentra indicating that such a  barrier layer was not forming. As implied from Fig. 4, a  visibly distinct silica-rich liquid/glassy surface layer was  not necessary for  a decelerating oxidation rate;  a de creasing mass change slope with time was observed for  the ZBS specimen with 37.1 mol% SiC, but  a glassy  layer was first observed for the ZBS specimen with 52.1  mol% SiC. The formation of enough silica to seal around  zirconia particles near  the surface would function as a  diffusion barrier, even if a distinct amorphous  surface  layer is not visible.  SiC concentrations in the ZBS series at and above 44.4  mol% are impressively effective at forming a thin passi vating layer. Under  rapid impact with the atmosphere,  as would be seen in a leading-edge aerospace applica tion,  such a high concentration of SiC may result  in a  liquid/glassy layer which may rapidly ablate  away.  In  contrast, a lower concentration of SiC might be expected  to form an amorphous surface layer with anchored zirco nia crystalline phases  interdispersed, providing protec tion against  the silica being convected away by a high  velocity  gas  stream. This  assumes  that  the  dominant  mechanism of glassy phase removal  is not vaporization  from frictional heating.  An interesting result  from the specimens oxidized at  1500  \\x0e  C is that  the silica layer thickness decreased, and  oxidation resistance increased, with increasing SiC con tent. For compositions of  low SiC content,  since there  would be an inadequate concentration of silica formed to  fully infiltrate the extensive amount of zirconia formed,  a seal would not  initially develop, and further oxidation  penetrating into the specimen would continue. Eventual ly, enough silica would form so that it could flow toward  the surface, filling in the space between zirconia crystals  with liquid, forming a coherent sealing layer, appearing  as the glassy surface layer  in the cross-section electron  micrographs. For higher SiC content  specimens, more  liquid phase is  formed relative to zirconia, allowing a  sealing layer  to develop with less zirconia surface area  to infiltrate, and such glassy layers would be less thick.  The oxide underlayers would correspondingly also be  less thick, as less SiC must be oxidized to supply liquid  phase to form the sealing layer.  The question remains as  to why the liquid phase is  drawn to near-surface regions to form the distinct glassy  surface layer observed, leaving a porous region of zirco nia,  deficient  in silica,  behind: A completely wetting  liquid will  flow to coat all solid surfaces, while a non wetting  liquid  tends  to  bead  up  into  a  sphere,  being  self-cohesive and minimizing contact with the solid as  FIG. 7. TEM micrograph of the amorphous surface layer of ZBSTB5  heat-treated at 1500  \\x0e  C, along with EDS spectra of selected regions.  Copper EDS peaks  correspond to the  copper grid upon which the  specimen was mounted, and gallium peaks correspond to embedded  gallium FIB gas milling atoms.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1862  \\x0c', 'F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  FIG. 8. TEM micrographs of the oxide underlayer of ZBSTB5 (16.61 mol% TaB2) heat-treated at 1500  \\x0e  C, along with EDS spectra of selected  regions.  FIG. 9. Effects of heat  treatment  in the TG in flowing air  for compositions of ZrB2-B4C-SiC, with varying concentrations of TaB2 (ZBSTB  series).  (a) Oxide underlayer  thicknesses were estimated from SEM micrographs of specimen cross section fracture surfaces. Solid and dashed  lines, and indicated temperatures, have the same meaning as in Fig. 4. (b) Mass change after soaking at  the indicated temperatures for 4 h in the  TG. Values corresponding to 0 mol% TaB2 were from sample ZBS10.  well as minimizing its own surface area. Silicate liquids are known to be non-wetting to zirconia.21 This is con remains  adequately wetting  to  zirconia  to  form an  even surface coating,  rather  than forming beads on the  sistent with the newly-formed borosilicate liquid phase  surface. The wicking direction may also be a result of  in  this work  being wicked up  to  the  initially-formed  the newly-formed liquid phase having a lower viscosity  liquid phase at  the surface. The liquid phase apparently  than the liquid phase at  the surface, since the latter may  http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  J. Mater. Res., Vol. 24, No. 5, May 2009  1863  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  be more silica-rich from B2O3 volatilization. The more  fluid liquid would more easily draw to the more viscous  liquid.  X-ray diffraction results (Table III) show that TaB2 or  TaSi2 additions to batch compositions form a ZrB2-TaB2  solid solution after sintering heat  treatments. In the case  of TaSi2 additions (ZBSTS), an enhanced SiC content  is  TaSi2(s) + 2B4C(s) ! TaB2(s) + 2SiC(s) + 6B(s) detected in XRD patterns. This implies the 1927 C = -111.0 kJ/mol) occurred, in which the TaB2  reaction: (DG  \\x0e  rxn,  \\x0e  product  formed a solid solution with ZrB2, and the B  product may be amorphous or below XRD detection.  During oxidation heat treatments, SiC oxidized to form  CO2(g) (initially) or CO(g) (subsequently) and amorphous  silica—the latter not being detectable via XRD. Phases  such as SiC and ZrB2 detected from XRD analysis of the  oxidized surfaces of  the ZBS series are interpreted as  being from phases  residing beneath the surface glassy  layers and the oxide underlayers. With TaB2 additions, in our previous work,20 it was shown that the TaC (again  detected in this work),  is a thermodynamically possible  reaction product of oxidation of TaB2 and SiC. TEM EDS  (Fig. 8) may imply slight solubility of tantalum in ZrO2,  though there is no indication of this from XRD results.  Increasing TaB2 additions to a ZrB2-SiC-B4C compo sition showed general,  though modest,  improvement  in  oxidation resistance at 1200 and 1400  \\x0e  C. At 1500  \\x0e  C,  small  (3.32 mol%) TaB2 additions  improved oxidation  resistance as well. These additions resulted in decreases  in the thicknesses of  the oxide underlayers, but had no  significant effect on the thicknesses of  the amorphous  surface  layers. This  improved  oxidation  resistance  is  interpreted to correspond to better  sealing in the oxide  underlayer. This may be from the greater wetting of the  liquid phase  to the TaC particles  formed (implied by  the right-hand micrograph in Fig. 8). Alternatively, and  the preferred explanation, is that the zirconia particle size  is substantially smaller (for equal amounts of SiC) when  TaB2 was introduced (Figs. 13 and 14). This would be  expected as oxidation of  the tantalum-zirconium boride  solid solution to produce segregated phases  (ZrO2 and  TaC) would result  in smaller particle sizes than the oxi dation of pure ZrB2 to form ZrO2 alone. The borosilicate  liquid phase which forms would have a greater tendency  to be entrapped in the oxide underlayers of finer particle  sizes. This would make  the  liquid-encapsulated oxide  underlayer more  impermeable  to atmospheric oxygen,  and increase overall oxidation resistance.  FIG. 10. TG traces of specimens of varying TaSi2 content oxidized in  air at  three different  temperatures. Dashed curves are 1200  \\x0e  C, solid  curves are 1400  \\x0e  C, and dot-dashed curves are 1500  \\x0e  C.  FIG. 11. Effects of heat treatment in the TG under flowing air for ZrB2-B4C-SiC with varying concentrations of TaSi2 (ZBSTS series). (a) Oxide  layer thicknesses estimated from SEM micrographs of specimen fracture surface cross sections. Solid and dashed lines and indicated temperatures  have the same meaning as in Fig. 4. (b) Mass changes after soaking at the indicated temperatures for 4 h in the TG. Values corresponding to 0 mol%  TaSi2 were from sample ZBS10.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1864  \\x0c', 'F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  FIG. 12. TG traces of specimens of varying TaB2/TaSi2 ratio (ZTBS series) oxidized in air at  three different  temperatures. Dashed curves are  \\x0e  \\x0e  1200  C, solid curves are 1400  C, and dot-dashed curves are 1500  C.  \\x0e  Remarkably, additions of TaB2 in excess of 3.32 mol%  resulted in linear mass  increases with time, or  in fact,  showed acceleration inflection points during the 4 h soak  periods (Fig. 5). An overall reduction in oxidation resis tance is seen with increasing TaB2 concentration (Fig. 9).  As the TaB2 content was increased, both amorphous sur face layer and oxide underlayer  thicknesses coarsened.  As can be seen in Fig. 14 (lower micrograph), zirconia  dendrites appear  to be growing through the amorphous  surface coating. The rectangular zirconia particle floating  in the glass  in Fig. 7 is  likely a broken-off portion of  a dendite, fractured away by convection currents in the liquid phase.12 It is speculated that the presence of tanta lum, perhaps slightly soluble in the glass, facilitated en hanced solution and precipitation of zirconia in the glassy  phase. After a time delay associated with a certain densi ty of these dendrites having grown to span from the oxide  underlayer to the glass surface, a discontinuous accelera tion in overall oxidation rate was induced. The dentrites  would cause  this by either  functioning as  channels of  fast oxygen anion conductors, or via imperfect wetting  (sealing) between liquid and the dendrites,  facilitating  enhanced oxygen penetration through interfacial fissures.  TaB2 is a more effective additive than TaSi2 at 1500  C,  when added in low concentrations (3.32 mol%) (see Figs. 9  and 11). The accelerating rate of oxidation at 1500  C with  TaB2 additions converts to a decelerating rate as TaB2 is  substituted with TaSi2. As with TaB2, additions of TaSi2  above  3.32 mol% TaSi2  reduces  oxidation  resistance,  though not  to nearly the extent observed with increasing  TaB2. Oxidation of  the  silicon provided  by  the TaSi2  additive increased the amount of  liquid phase relative to  zirconia, and changed the composition of the liquid phase  to one  that  is more  silica  rich. This  liquid phase may  be interpreted as being more  impermeable to diffusing  gaseous oxygen, and/or more inhospitable to solution and  precipitation of zirconia dendrites.  \\x0e  \\x0e  FIG. 13. Comparison of oxide underlayers at  the same magnification  for ZBS10, ZBSTB5, and ZBSTS4, oxidized at 1400  C.  \\x0e  http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  J. Mater. Res., Vol. 24, No. 5, May 2009  1865  \\x0c', 'http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  V. CONCLUSIONS  In the ZBS series, ZrB2 and SiC oxidized to form a  borosilicate  liquid  and ZrO2.  If  the  amount  of  liquid  phase was  inadequate,  it was  interpreted  that  newly formed liquid was wicked up to join liquid previously  formed at the specimen surface. Lower concentrations of  SiC resulted in thicker glassy surface layers, with high  concentrations of embedded zirconia, and thicker zirco nium oxide underlayers. Fixing the SiC content at 27.91  mol%, small (3.32 mol%) additions of TaB2 were shown  to  be  beneficial  in  improving  oxidation  resistance  at  1500  \\x0e  C. This was attributed to the oxidation and break up of TaB2-ZrB2 boride solid solution grains into fine scale ZrO2 and TaC particles, which better  retained an  encapsulating  liquid  phase.  Increasing  TaB2  content  beyond  3.32 mol% improved  oxidation  resistance  at  1200 and 1400  \\x0e  C; however, oxidation resistance was  precipitously  degraded  at  1500  \\x0e  C. The  presence  of  higher concentrations of  tantalum in some way facili tated the solution and precipitation of zirconia dendrites.  After  some  period  of  time,  these  dendrites  permitted  the acceleration of oxidation via providing a fast diffu sion path either through the dendrites or along the glass/  crystal  interfaces of  these dendrites. TaSi2  concentra tions above 3.32 mol% also resulted in diminished oxi dation  resistance,  but  on  a much-reduced  scale. The  higher silica content of  the borosilicate glass was inter preted to be less facilitating of the zirconia solution and  precipitation mechanism.  ACKNOWLEDGMENTS  This project was  funded by the Air Force Office of  Scientific Research, Contract FA9550-04-1-0140. The  authors would like to express their appreciation for  the  helpful suggestions and support of  their contract moni tor, Dr. Joan Fuller.  REFERENCES  1. Y. Murata and E.B. Whitney: Densification and wear  resistance  of ceramic systems:  III. Tantalum mononitride-zirconium dibor ide. Am. Ceram. Soc. Bull. 48(7), 698 (1969).  2. Y. Murata: Densification and wear resistance of HfN-ZrB2 compositions. Am. Ceram. Soc. Bull. 52(3), 255 (1973).  3. X. Zhang, P. Hu, S. Meng, J. Han, and B. Wang: Microstructure  and mechanical properties Mater. 312, 287 (2006).  of ZrB2-based  ceramics. Key Eng.  4. W.C. Tripp: Effect of an SiC addition on the oxidation of ZrB2.  Am. Ceram. Soc. Bull. 52(8), 1606 (1973).  5. M.M. Opeka,  I.G.  Talmy,  E.J. Wuchina,  J.A.  Zaykosi,  and  S.J. Causey: Mechanical,  thermal,  and oxidation  properties of  refractory hafnium and zirconium compounds. Soc. 19, 2405 (1999).  J. Eur. Ceram.  6. A.R. Rezaie, W.G. Fahrenholtz, and G.E. Hilmas: Oxidation of  zirconium diboride-silicon carbide at 1500 C in a low partial pressure of oxygen. J. Am. Ceram. Soc. 89(10), 3240 (2006).  \\x0e  7. A.K. Varshneya: Fundamentals of Inorganic Glasses (Academic  Press, New York, 1994).  8. E.J. Opila and M.C. Halbig: Oxidation of ZrB2-SiC. Elec. Chem.  Soc. Proc. 12, 221 (2002).  9. A. Rezaie, W.G. Fahrenholtz,  and G.E. Hilmas: Evolution  of  structure during the oxidation of zirconium diboride-silicon car bide in air up to 1500  \\x0e  C. J. Eur. Ceram. Soc. 27, 2495 (2007).  10. M. Opeka, I. Talmy, and J. Zayk:oski: Oxidation-based materials  selection for 2000  \\x0e  C+ hypersonic aerosurfaces: Theoretical con siderations  and  historical  experience.  J. Mater.  Sci.  39,  5887  (2004).  11.  J. Han, P. Hu, X. Zhang, and S. Meng: Oxidation behavior of  zirconium diboride-silicon carbide at 1800  \\x0e  C. Scr. Mater. 57,  825 (2007).  12. S.N. Karlsdottir, J.W. Halloran, and A.N. Grundy: Zirconia trans port by liquid convection during oxidation of zirconium diboridesilicon carbide. J. Am. Ceram. Soc. 91(1), 272 (2008).  13.  I.G. Talmy, J.A. Zaykoski, M.M. Opeka, and S. Dallek: Oxida tion of ZrB2 ceramics modified with SiC and group IV-VI transition metal diborides. Elec. Chem. Soc. Proc. 12, 144 (2001).  14. W. Vogel: Glass Chemistry, 2nd ed. (Springer-Verlag, New York,  1994).  15. E. Opila, S. Levine, and J. Lorincz: Oxidation of ZrB2and HfB2 based ultra-high temperature  ceramics: Effect of Ta  additions.  J. Mater. Sci. 39, 5969 (2004).  FIG. 14. Comparison of oxide underlayers at  the same magnification  for ZBS10, ZBSTB5, and ZBSTS4, oxidized at 1500  \\x0e  C.  F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  J. Mater. Res., Vol. 24, No. 5, May 2009  1866  \\x0c', 'F. Peng et al.: Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride  16.  I.G. Talmy, J.A. Zaykoski, M.M. Opeka, and A.H. Smith: Properties of ceramics in the system ZrB2-Ta5Si3. J. Mater. Res. 21(10),  2593 (2006).  19. X-H. Zhang, P. Hu, J-C. Han, L. Xu and S-H. Meng: The addition  of lanthanum hexaboride to zirconium diboride for oxidation resistance. Scr. Mater. 57(1) 36-39 (2007).  improved  17. S. R. Levine, and E.  J. Opila: Tantalum addition to zirconium  diboride  for  improved  oxidation  resistance. NASA/TM-2003-  20. F. Peng and R.F. Speyer: Oxidation resistance of fully dense ZrB2 with SiC, TaB2, and TaSi2 additives. J. Am. Ceram. Soc. 91(5),  212483.  1489 (2008).  18. S.C. Zhang, G.E. Hilmas, and W.G. Fahrenholtz:  Improved oxi 21. W.E.  Lee  and W.M.  Rainforth:  Ceramic Microstructures,  dation resistance of zirconium diboride by tungsten carbide additions. J. Am. Ceram. Soc. 91, 3530 (2008).  Property Control by Processing (Chapman and Hall, London,  1994).  http://journals.cambridge.org  Downloaded: 18 Mar 2015  IP address: 140.182.176.13  J. Mater. Res., Vol. 24, No. 5, May 2009  1867  \\x0c']"
},{
  "_id": 54,
  "PDF": "Effect of surface oxidation on thermal shock resistance of ZrB 2 –SiC–ZrC ceramic at temperature difference from 800 to 1900 °C.pdf",
  "Text": "['Corrosion Science 98 (2015) 233-239  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effect of surface oxidation on thermal shock resistance of ZrB2-SiC-ZrC ceramic at temperature difference from 800 to 1900  C  Zhi Wang a,∗ , Peng Zhou b , Zhanjun Wu a  a School of Aeronautics and Astronautics, Faculty of Vehicle Engineering and Mechanics, State Key Laboratory of Structural Analysis for Industrial  Equipment, Dalian University of Technology, Dalian 116024, China b School of Natural Sciences and Humanities, Harbin Institute of Technology Shenzhen Graduate School, Shenzhen 518055, China  a r t i c l e  i n f o  a b s t r a c t  1900   C through the In the present work, the ZrB2 -ZrC-SiC ceramic was fabricated at reaction (5Zr + 2B4 C + Si → 4ZrB2 + ZrC + SiC). The effect of oxidation behavior on the thermal shock resistance of the ZrB2 -ZrC-SiC ceramic was investigated in two kinds of heat conditions (air and vacuum) by measuring the residual strength after water quenching for the temperature difference ranging from 200 to 1900   C. The improvement of thermal shock resistance was attributed to the formation of oxide layers. The results of the present work indicate that the short-term oxidation is favorable to improve the thermal shock resistance of ZrB2 -based ceramics.  © 2015 Elsevier Ltd. All rights reserved.  Article history:  Received 12 March 2015 Accepted 15 May 2015 Available online 31 May 2015  Keywords:  A. Ceramic B. SEM B. XRD C. Oxidation  1.  Introduction  Known as ultrahigh temperature ceramics [1], transition metal diborides, carbides, and nitrides show a unique combination of properties such as high melting temperature, chemically inert, high strength, high thermal and electrical conductivity, and extremely hard solids with high thermal stabilities [2,3]. As a result, these ceramics are potential candidates for a variety of high temperature structural applications [4], such as leading edges and nose caps in hypersonic re-entry space vehicles, rocket nozzle inserts and air-augmented propulsion system components in hypersonic aerospace vehicles where materials can be subjected to temperatures between 1600  C and 2800  C in an oxidizing environment of high velocity dissociated air [5]. Zirconium diboride (ZrB2 ) has received especially strong attention because of it has a relatively low density compared with the other candidates for ultrahigh temperature applications [6]. In recent years, the development of the ZrB2 -based ceramics is driven by the demand of developing re-usable hot structures as thermal protection systems of space vehicles able to re-enter from low earth orbit at relatively high speed (order of 8 km/s) [7]. External thermal protection systems of high speed vehicles in the form of rigid surfaces in areas of ultra-high oper(>1800  C) must be ating temperature able to withstand high  ∗ Corresponding author. Tel.: +86 411 84706791; fax: +86 411 84706791. E-mail address: wangzhi1780@126.com (Z. Wang).  http://dx.doi.org/10.1016/j.corsci.2015.05.035 0010-938X/© 2015 Elsevier Ltd. All rights reserved.  temperature, high heat flux and mechanical strength associated with vibrations at launch and re-entry in Earth’s atmosphere [8-11]. Therefore, it is necessary to develop the thermal protection materials endowed with good resistance to oxidation and thermal shock. The oxidation induced by high temperature is a major issue for the use of the ZrB2 -based ceramics [9]. In the past 10 years, groups in the United States, Japan, Italy and China have investigated the oxidation behavior of ZrB2 -based ceramics [12-15]. Pure ZrB2 forms a protective zirconia (ZrO2 ) scale under oxidizing environments, but the vaporization of gaseous oxidation product e.g., B2O3 can increase the porosity of the surface oxide layer and reduce the oxidation resistance of this materials [15-17]. However, recently important work has shown that ZrB2 -SiC ceramics have significantly improved oxidative resistance compared with pure ZrB2 . Fahrenholtz et al. [18] have studied the microstructure of oxide layers formed on ZrB2 -SiC ceramics after oxidation at temperatures up to 1500  C. They have indicated that the typical layer composed of three layers: (1) a SiO2 -rich glassy layer; (2) a thin ZrO2 -SiO2 layer; (3) a SiC-depleted layer. The development of the layered structure has analyzed with the aid of a thermodynamic model that involved volatility diagrams for ZrB2 and SiC [19]. When exposed to air at high temperatures of below 1800  C, ZrB2 reacts to form crystalline ZrO2 and boria (B2O3 ) which is a quite fluid liquid. The SiC reacts to form silica (SiO2 ) which is a very viscous liquid. Typically, a twolayer oxide film forms, with a zirconium oxide inner film and a SiO2 outer film. The SiO2 can dissolve with the B2O3 to form a B2O3 -SiO2 liquid solution of intermediate viscosity. At the temperature of oxidation, B2O3 is much more volatile than SiO2 . At the outer surface of  \\x0c', '234  Z. Wang et al. / Corrosion Science 98 (2015) 233-239  the oxide, the B2O3 is preferentially evaporated, so that the liquid film at the outer surface is predominantly viscous SiO2 . The microstructural evidence suggests that oxygen is transported inward through the oxide film, to react at the interface between the oxide and the ZrB2 [20]. However, the oxidation resistance of ZrB2 -SiC ceramics markedly degrades beyond 1800  C, as the SiC actively oxidizes to form gaseous SiO and CO or the SiO2 layer volatilizes [21]. Additionally, there is evidence that incorporating ZrC into the ZrB2 -SiC ceramics gives even greater protection against oxidation due to increasing the content of ZrO2 in oxide layer [22]. Furthermore, in machining process of ZrB2 -based ceramics, whether electron discharge machining or a variety of precision grinding equipment is used, which inevitably results in the presence of the surface defects such as cracks, holes and so on [23]. ZrB2 -based ceramics, because of inherently low fracture toughness, are susceptible to catastrophic fracture caused by the thermal stresses under the thermal shock conditions [24]. When the ZrB2 based ceramics are loaded by the thermal stresses, the defect-tip will produce stress concentration. Concentration of stress can be calculated according to Griffith formula [25]:  (cid:2)  (cid:3)  (cid:4)  (cid:2)m =  1 + 2  C  (cid:3)  (cid:2)  (1)  where (cid:2)m is concentrated stress, (cid:2) is average strength of materials, C is size of defect and (cid:3) is radius of defect-tip. The ZrB2 -based ceramics are thus destroyed at low stress because the radius of defect-tip is much less than the size of defect. Recently, many investigators have confirmed that the ceramic matrix composites containing Si, Ti and Al, such as Al2O3 /SiC, Si3N4 /SiC, Ti3AlC2 , have the ability to heal surface defects so as to increase the thermal shock resistance through surface oxidation treatment [26-28]. Although great efforts have been separately devoted to the oxidation behavior and thermal shock of ZrB2 -based ceramics and much progress has been gained recently, up to date, there are only a few papers devoted to the effect of oxidation behavior on thermal shock resistance of ZrB2 -based ceramics, especially the oxidation in short time. In the present work, ZrB2 -15.05vol%ZrC-12.15vol%SiC (ZrB2 -ZrC-SiC) ceramic with raw powders of Zr-B4C-Si through the reaction (2) was fabricated at 1900  C for 1 h under a uniaxial load of 30 MPa in Ar atmosphere. 5Zr + 2B4C + Si → 4ZrB2 + ZrC + SiC  (2)  The effect of oxidation behavior on thermal shock resistance of the ZrB2 -ZrC-SiC ceramics was investigated in two kinds of heat conditions (air and vacuum) by measuring the retention of the flexural strength after water quenching for the temperature difference ranging from 200 up to 1900  C. The purpose of this paper is to report the effect of oxidation behavior on thermal shock resistance of ZrB2 -based ceramics, which can be applied to aid materials engineering design for the development of quality assurance and characterization assessment of durability.  2. Experimental procedures  Commercially available Zr (mean particle size 30.0 \\u242em, Northwest Institute for Non-ferrous Metal Research, China), B4C (mean particle size 5.0 \\u242em, Jingangzuan Boron Carbide Co. Ltd., Mudanjiang, China) and Si powders (mean particle size 20.0 \\u242em, Shouchang Co. Ltd., Shanghai, China) were used as raw materials. The raw Zr-B4C-Si powders in molar ratio 5/2/1 according to reaction (2) were ball-milled in ethanol for 12 h in a nylon-coated tank using ZrO2 balls. Rotary evaporation was used to separate the milled powder from the slurry. Bulk ZrB2 -ZrC-SiC ceramic was fabricated by in situ hot-pressing reaction synthesis. The milled powder was heated at 10  C/min from room temperature  to 1500  C and then held at this temperature for 60 min in argon atmosphere. After the hold, the milled powder was heated at 10  C/min to 1900  C and held for another 60 min, the application of pressure began at 1800  C and gradually increased to 30 MPa before the temperature reached a value of 1900  C. The bulk density of the ZrB2 -ZrC-SiC ceramic was measured using the Archimedes’ technique with deionized water as the immersing medium. The relative density was determined by dividing the bulk density by the theoretical density based on a rule of mixture calculation. The phase composition was determined by X-ray diffraction (XRD; Rigaku, Dmax-rb, CuK␣ = 1.5418 ˚A). The sample was quenched from various elevated temperatures into water and determining the temperature difference that resulted in a reduction of strength for In the present work, a water bath of 20  C a rectangular sample. was used to measure experimentally the critical temperature difference which was to evaluate the thermal shock resistance. The critical temperature difference is defined as 70% of the room temperature strength, which was determined using linear interpolation of the retained strength values as described in ASTM C1525-04 [29]. The polished rectangular bars for thermal shock testing were heated in two kinds of heat conditions (air and vacuum) up to the desired temperature difference (the temperature difference is defined as the difference between the sample temperature and water temperature, for example, the temperature difference of 1000  C (1020 − 20): the sample is heated at 1020  C, the temperature of water bath is 20  C) and held for 10 min to eliminate any temperature gradient effect before quenching by dropping parallel to their tensile surface into the water bath. The time taken for the transfer from the furnace to the water bath was less than 1 s. The temperature differences for water quenching were 200, 300, 400, 500, 600, 700, 800, 900, 1000, 1100, 1200, 1300, 1400, 1500, 1600, 1700, 1800, 1900  C, respectively. Flexural strength of the samples before and after the water quenching was tested in three-point bending on 3 mm × 4 mm × 36 mm bars, using a 30 mm span and a crosshead speed of 0.5 mm min−1 . Before water quenching, all samples were ground and polished with diamond slurries down to a 1 \\u242em finish and the edges of all samples were chamfered to minimize the effect of stress concentration due to machining defects. At least six samples were tested for each experimental condition and all samples were from same billet. The microstructural feature of the sample was observed by scanning electron microscopy (SEM, FEI Sirion, Holland) along with energy dispersive spectroscopy (EDS, EDAX Inc.) for chemical analysis.  3. Results and discussion  3.1. Microstructure  An XRD spectrum obtained from the polished surface of the hot-pressed ZrB2 -ZrC-SiC ceramic is shown in Fig. 1A. Apparently, the phase analysis indicated the predominant phases for the assintered ceramic were ZrB2 , SiC and ZrC, no peaks of inpurities were detected, which revealed that the reaction of Zr-B4C-Si powder is complecated after the reactive hot-pressing. The ZrB2 -ZrC-SiC ceramic used to produce bars for flexural strength test had bulk densities of 5.80 g/cm3 . Using a rule of mixture calculation and assuming that the true densities are 6.09 g/cm3 for ZrB2 , 3.21 g/cm3 for SiC and 6.73 g/cm3 for ZrC, the theoretical density of the ZrB2 -15.05vol%ZrC-12.15vol%SiC ceramic was calculated to be 5.83 g/cm3 . Based on this true density, the relative density of the ZrB2 -ZrC-SiC ceramic was 99.49%. Fig. 1B shows the polished surface of ZrB2 -ZrC-SiC ceramic. Based on EDS spectra (not shown here), it could be seen that the microstructure of ZrB2 -ZrC-SiC ceramic was characterized by the gray ZrB2 , dark SiC as well as the light ZrC phases. Furthermore, the present of the cracks and  \\x0c', 'Z. Wang et al. / Corrosion Science 98 (2015) 233-239  235  Fig. 1. XRD spectra (A) obtained from the polished surface of the hot-pressed ZrB2 -ZrC-SiC ceramic and the micrograph (B) of the ZrB2 -SiC-ZrC ceramic.  holes on the polished surface of ZrB2 -ZrC-SiC ceramic were due to significant pullout of crystalline grains occurred during the polishing process, not due to incomplete densification because the relative density of the ZrB2 -ZrC-SiC ceramic was 99.49% [30]. Compared with 800 MPa of most ZrB2 -based ceramics, the flexural strength was slightly reduced to 580 MPa due to the greater grain size of SiC phase because flexural strength was a strong function of SiC grain size. Fig. 2 shows the residual strength for the samples heated in air and vacuum at different temperature differences. When the temperature difference is more than 300  C, the residual strength values for the samples heated in vacuum gradually decreased as the temperature difference increased up to 700  C. The variability in the residual strength values quenched at the temperature difference of 800  C was significant because the individual sample failed upon quenching [3,4]. Generally, the ZrB2 -based ceramics inevitably contain some defects in the form of porosity, microcracks, impurity and so on [23,29]. When the ZrB2 -based ceramics are subjected to the thermal stresses arising from the temperature gradients, the surface of the ZrB2 -based ceramics is placed under tension and the interior under compression [3,9,17]. A stress gradient is expected at the edges and corners of the flexural bars due to higher proportion of surface area in these regions, which results in more rapid cooling of these regions compared to the bulk of the sample [29]. If a critical defect exists near an edge or corner, a sample may thermally shock at lower temperature difference than samples with similar sized defects that are across the width of the tensile surface or in the bulk. The distribution of critical defects and stress gradients may also be the cause of the variability in the  Fig. 2. Residual strength of the ZrB2 -SiC-ZrC ceramic for two kinds of heating conditions as a function of thermal shock temperature difference.  strength for samples quenched at the temperature difference of 800  C. All of sample failed upon quenching when the temperature difference was more than 800  C. The large numbers of paper on investigating the thermal shock behavior of the ceramic matrix composites have published and it is recognized that the decay in flexural strength measured after the water quenching is theoretically attributed to the micro-defects appeared on the surface of the sample [3,4,9,17,23,29]. Unfortunately, no defects on the surface of the sample are experimentally observed in the published papers. In the present work, the macroscopic cracks were observed on the surface of the sample, which experimentally confirmed that the decay in flexural strength after the water quenching was mainly attributed to the macroscopic cracks appeared on the surface of the sample, as show in Fig. 3. In contrast to the samples heated in vacuum, the residual strength values for the samples heated in air increased gradually as the temperature difference increased from 800 up to 1200  C; and the residual strength values rapidly increased from 1200 up to 1300  C, then slightly increased further from 1300 up to 1700  C and slightly decreased from 1700 up to 1900  C, as shown in Fig. 2. The residual strength values of the sample quenched at temperature differences of >1200  C were equivalent to original strength values of the ZrB2 -SiC-ZrC ceramic. The enhancement in the residual strength for the samples heated in air was contradictory to the prediction based on the conventional damage mechanics [25,31], which more indicated the effect of the microstructure evolution due to the surface oxidation on residual strength because the SEM observations of the fractured surfaces of all samples did not show the significant differences compared with one of the unquenched samples. Furthermore, the critical temperature difference to represent thermal shock resistance of ceramic matrix composites can be measured experimentally by ASTM C1525-04. Although the residual strength values for the samples heated in air are obviously different from that for the samples heated in vacuum at temperature differences from 200  C to 800  C, the measured \\x01Tcrit values for the sample heated in air and vacuum are statistically identical at 362-356  C. SEM analysis has confirmed no change of the surface microstructure of sample quenched in air and vacuum at temperature difference of <500  C, compared with unquenched sample. The surface macrographs (not shown here) for the samples heated in vacuum are similar to one for the samples heated in air at the temperature difference ranging from 200 to 500  C. Compared with the unquenched samples, the surface of the sample darkened when the temperature difference was higher 500  C due to the slight surface oxidation of the sample. For example, ZrB2 is oxidized to ZrO2 and B2O3 at above 650  C [32]; SiC is oxidized to SiO2 and CO2 at above 900  C [18]; ZrC is oxidized to ZrO2 and CO2 at above  \\x0c', '236  Z. Wang et al. / Corrosion Science 98 (2015) 233-239  Fig. 3. The surface macrograph for the quenched samples (3 mm × 4 mm × 36 mm) at temperature differences of 800   C (B) before the water quenching (A).  450  C [33]. The macrographs for the samples quenched at 800  C in vacuum are shown in Fig. 3 and the obvious cracks are readily detected, compared with the sample before water quenching. This was attributed mainly to that thermal stress was greater than the original strength of the ZrB2 -SiC-ZrC ceramic. ZrC(s) + 2O2 → ZrO2 (s) + CO2 (g) ZrB2 (s) + 5/2O2 → ZrO2 (s) + B2O3 SiC(s) + 2O2 → SiO2 (l) + CO2 (g)  (4)  (3)  (5)  The slight surface oxidation of the ZrB2 -SiC-ZrC ceramic was not detected by SEM due to short holding time when the temperature difference was below 800  C. The surface oxidation can  be observed on the surface of the samples quenched at the temperature difference of ≥800  C. ZrC grains were oxidized to CO2 and ZrO2 , and ZrB2 grains began to oxidize to B2O3 and ZrO2 identified by Based on EDS spectra (not shown here), as shown in Fig. 4, and the obvious microdefects such as cracks and pits that appeared at interfaces between ZrB2 (ZrC) and ZrO2 grains were attributed to both the thermal stress during water quenching and the volume change during the reaction from ZrB2 (ZrC) to ZrO2 (300% volume expansion based on density calculations) [18]. When the temperature difference was 800  C, some ZrB2 (ZrC) grains in the surface of the samples were oxidized to ZrO2 as shown in Fig. 4B. An amount of ZrO2 grains increased as the temperature difference increased up to 1100  C, as shown in Fig. 4. Based on a negative standard Gibbs  Fig. 4. SEM images of the surface of sample quenched at temperature difference of 800 (A and B) and 1100 (C and D)   C for heated in air; low magnification (A and C) and high magnification (B and D).  \\x0c', 'Z. Wang et al. / Corrosion Science 98 (2015) 233-239  237  Fig. 5. SEM images of the surface of sample quenched at temperature difference of 1200 (A), 1300 (B), 1400 (C), and 1500 (D)  C for heated in air.  free energy of reaction, B2O3 can react spontaneously with water during water quenching according to Eq. (6) [34], resulting in the that the expected B2O3 glass is not observed on the surface of the samples. B2O3 + 3H2O → 2H3BO3 \\x01G0  = −28.84 kJ/mol  298  (6)  The SiC grains are not oxidized due to the holding time limit because the oxidation of SiC is much slow at below 1200  C [18]. Although the SiO2 can react with B2O3 to produce stable SiO2 ·B2O3 glass that can be eutectic at 372  C and shows better water resistance, reduced volatility at high temperature, and higher viscosity the SiO2 ·B2O3 glass does not than B2O3 glass. However, form in water quenching conditions, which is presumably attributed to that the oxidation of SiC below 1200  C is limited by the holding time and high heating temperature leads to B2O3 rapid evaporation, especially when the temperature approaches and exceeds 1100  C. The amount of surface defects increased as the temperature difference increased from 800  C to 1100  C, however, the residual strength values also increased as the temperature difference increased. This phenomenon may also be the cause of the appearance of the compressive stress resulted from volume expansion based on density calculations due to the conversion of ZrB2 (ZrC) and SiC to ZrO2 and SiO2 . The appearance of the compressive stress zone beneath the surface oxide layers is favorable to inhibit the crack initiation and propagation to damage matrix of the ZrB2 -SiC-ZrC ceramic during the transient water quenching, which is beneficial to resistance to thermal shock because the thermal stress is consumed partially due to the appearance of the compressive stress. The temperature difference was 1200  C and the surface of sample was coated by thinner oxide layer, as shown in Fig. 5A. B2O3 was much more volatile than SiO2 . At the outer surface of the oxide layer, the B2O3 was preferentially evaporated, so that the liquid  film at the outer surface was predominantly viscous SiO2 . Although quantitative analysis is not possible by EDS since B is at the limit of the detection capability or by XRD since the B2O3 is amorphous, some B2O3 inevitably remains in the SiO2 in this layer. The small quantity of liquid B2O3 -SiO2 solution of intermediate viscosity can heal the defects in the form of microcracks and pits on the surface of the samples, which is favorable to the increase in residual strength of the ZrB2 -SiC-ZrC ceramic. increased up to 1300  C, When the temperature difference as shown in Fig. 5B, the surface of sample was coated by B2O3 -SiO2 -ZrO2 oxide layer, which had the appearance of islands of ZrO2 particles clustered on the outer surface of the oxide film. The ZrO2 islands were surrounded by a region of SiO2 -rich glass, which resemble a lagoon surrounding the ZrO2 islands. Within the SiO2 -rich lagoons, and adjacent to the ZrO2 islands were lobes of B2O3 -rich glass arranged like the petals of a flower. These features can be interpreted as convection cells arising from flow of a liquid oxide solution containing B2O3 , SiO2 and ZrO2 , and the convection cells can effectively heal the surface defects because of the mutual wetting behavior of the different materials [35]. As the temperature difference further increased, the amount of SiO2 in outer oxide layer obviously increased whereas the amount of zirconia decreased, as shown in Fig. 5C and D, such a structure is favorable to the residual strength of the ZrB2 -SiC-ZrC ceramic after water quenching. For the sample quenched at the temperature difference of 1600  C, the surface of sample was coated by SiO2 -rich layer and a small quantity of the ZrO2 islands were still detected, as shown in Fig. 6A. However, the amount of borosilicate in outer oxide layer obviously decreased due to the evaporation of borosilicate at above 1700  C, as shown in Fig. 6B, and thin borosilicate glass layer fissures (black arrowhead) due to volume shrinkage under water quenching for the sample quenched from temperature difference of  \\x0c', '238  Z. Wang et al. / Corrosion Science 98 (2015) 233-239  Fig. 6. SEM images of the surface of sample quenched at temperature difference of 1600 (A), 1700 (B), 1800 (C), and 1900 (D)  C for heated in air; the inset is high magnification of 1900   C.  1800  C. Furthermore, the defects (white arrowhead) resulted from the escape of the gaseous products through the borosilicate layer were readily observed on the surface of the sample, as shown in Fig. 6C. When the sample was heated at the temperature of 1920  C (temperature difference of 1900  C), the borosilicate glass layer on the surface of sample entirety escaped, which resulted in that the surface of sample was coated by ZrO2 layer, as shown in Fig. 6D. The defects inevitably appear on the surface of the sample during the machining process because ZrB2 -SiC-ZrC ceramic is brittle materials. When the sample was quenched into the water, the thermal stress concentration occurred at defect-tip. (I) The defects on the surface of the sample were healed, which could reduce the thermal stress concentrated at defect-tip so as to increase the thermal shock resistance based on Eq. (1) described in Section 1. (II) The literature [31] has reported stress gradients with a maximum tensile stress at the end of the bar (edges) and at the surface of the mid-line of the width of the bar. Therefore, the appearance of the compressive stress zone beneath the surface oxide layers is favorable to inhibit the crack initiation and propagation during the transient water quenching prior to final failure, which is beneficial to resistance to thermal shock because the thermal stress is counteracted partially due to the appearance of the compressive stress. (III) Compared with the high thermal conductivity (90-120 W m−1 K−1 ) [36] of the ZrB2 -SiC-ZrC ceramic, the surface oxide layers of the low thermal conductivity (2-4 W m−1 K−1 ) [37,38] can act as the thermal barrier layer of the sample substrate. (IV) During the water quenching, the surface temperature of the oxide layers is reduced to the temperature of the water bath, whereas the surface of the sample substrate (close to interface between the oxide layer and the substrate) is still retained at higher temperature due to the low thermal conductivity of the oxide layers. The temperature difference between the surface of the sample substrate and the interior of the substrate is reduced due to the presence of the  oxide layers, which can result in the decrease in the thermal stress. (V) Furthermore, the thermal stress was consumed by the breaking of the oxide layers during the water quenching, thus protecting the sample substrate and enhancing the thermal shock resistance of the sample substrate. These positive factors due to the formation of oxide layer after heating at the temperature difference of >800  C make the improvement in residual strength of the ZrB2 -SiC-ZrC ceramic after water quenching. The results of the present work indicate short-term oxidation is favorable to improve the thermal shock resistance of ZrB2 -based ceramics.  4. Conclusions  In the present work, the effect of oxidation on thermal shock resistance of the ZrB2 -ZrC-SiC ceramics was investigated in two kinds of heat conditions (air and vacuum) by measuring the retention of the flexural strength after water quenching for the temperature difference ranging from 200 up to 1900  C. The results indicated that the improvement of thermal shock resistance was attributed to five factors due to the formation of oxide layers: (I) the healing of the surface defects, (II) the increase in the flexural strength, (III) the appearance of the compressive stress zone beneath the surface oxide layers, (IV) the decrease in the thermal stress, and (V) consumption of the thermal stress. The results here pointed to a promising method for improving strength and thermal shock resistance of ZrB2 -based ceramics.  Acknowledgements  This work was supported by the National Natural Science Foundation of China (51102031) and supported by the Fundamental Research Funds for the Central Universities DUT14LK26.  \\x0c', 'Z. Wang et al. / Corrosion Science 98 (2015) 233-239  239  References  [1] P. Hu, Z. Wang, Flexural strength and fracture behavior of ZrB2 -SiC ultrahigh temperature ceramic composites at 1800   C, J. Eur. Ceram. Soc. 30 (2010) 1021-1026. [2] Z. Wang, Z.J. Wu, G.D. Shi, The oxidation behaviors of a ZrB2 -SiC-ZrC ceramic, Solid State Sci. 13 (2011) 534-538. [3] S.B. Zhou, Z. Wang, X. Sun, J.C. Han, Microstructure, mechanical properties and thermal shock resistance of zirconium diboride containing silicon carbide ceramic toughened by carbon black, Mater. Chem. Phys. 122 (2010) 470-473. [4] Z. Wang, Q. Qu, Z.J. Wu, G.D. Shi, The thermal shock resistance of the ZrB2 -SiC-ZrC ceramic, Mater. Des. 32 (2011) 3499-3503. [5] M.M. Guron, M.J. Kim, L.G. 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  "_id": 55,
  "PDF": "Effect of tantalum addition on microstructure and oxidation of spark plasma sintered ZrB2-20vol_ SiC composites.pdf",
  "Text": "['Ceramics International 45 (2019) 13799-13808  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Effect of tantalum addition on microstructure and oxidation of spark plasma sintered ZrB2-20vol% SiC composites  T  Sravan Kumar Thimmappa, Brahma Raju Golla∗  Metallurgical and Materials Engineering Department, National  Institute of Technology, Warangal, 506 004,  India  A R T I C L E  I N F O  A B S T R A C T  Keywords: ZrB2 SiC Tantalum Spark plasma sintering Microstructure Oxidation  1.  Introduction  full density (> 99% theoretical density (ρth)) was achieved for ZrB2-20vol% SiC-Xwt.% Ta (X = 2,5, 5 Almost and 10) composites after Spark Plasma Sintering (SPS) (Temperature: 1900 °C, Pressure: 50 MPa; Time: 3 min). The microstructure of ZrB2-based composites exhibited core-rim structure and it consists of major crystalline phases (ZrB2 core, (Zr, Ta)B2 rim, SiC), minor amounts of ZrO2 and (Zr, Ta)C solid solution phases. Both the (401-195 μm) of ZrB2-20vol% SiC specific weight (from 22.91 to 18.77 mg/cm2) and oxide layer thickness composites decreased with increasing addition of Ta after the isothermal oxidation at 1500 °C for 10 h in air. The cross-sectional microstructure of oxidized samples displayed presence of a stack of three distinctive layers, which includes thick dense SiO2 top layer, SiC depleted intermediate layer and unreacted bulk. The present work clearly demonstrated the advantage of tantalum addition in improving the oxidation resistance of ZrB2-20vol% SiC.  Transition metal (Zr, Hf) borides, carbides and nitrides belong to the group of Ultra-High Temperature Ceramics (UHTCs) as they exhibit melting point above 3000 °C [1]. Among the family of transition metal boride UHTCs, zirconium diboride (ZrB2) is the most extensively studied ceramic as it exhibits good combination of properties: low density (6.08 g/cc), high melting point (3250 °C), high thermal conductivity (60-130 W/mK), high electrical conductivity and good elevated temperature properties [1-3]. Hence, ZrB2 is one of the key materials for the application of thermal protective systems and wing leading edges of hypersonic vehicles. The applications of phase pure ZrB2 is limited due to the processing difficulties, moderate fracture toughness and poor oxidation resistance at elevated temperatures [4-9]. Because of its strong covalent bonding nature and low self-diffusion coefficient, it is challenging to densify monolithic ZrB2 even at very high sintering temperatures. Different metallic or non-metallic reinforcements have been essentially added to enhance the densification of ZrB2. At low temperatures (i.e. up to 1100 °C), ZrB2 exhibits good oxidation resistance as it forms B2O3 layer at its surface during oxidation, which acts as a passive layer for the inward diffusion of oxygen. But at high temperatures, B2O3 will evaporate easily and the oxidation of ZrB2 progresses quickly [9]. Hence, ZrB2-based composites have been widely researched,  particularly, Si containing reinforcements such as SiC, MoSi2, TaSi2 etc. have been added to ZrB2 for improving its mechanical and thermal properties at elevated temperatures [10-15]. For examples, the addition of SiC enhanced the mechanical properties as it inhibited ZrB2 grain growth, improved thermal conductivity and the oxidation resistance of ZrB2 as well [10-17]. Recently, attempts were also made to study the effect of transition metals addition on the densification and oxidation resistance of ZrB2 [18,19]. Dehdashti et al. [19] revealed that Nb addition to ZrB2 resulted (Zr, Nb)B2 solid solution phase in the microstructure, and this phase reportedly improved the oxidation resistance of ZrB2 by forming Nb2Zr6O17. Silvestroni et al. [20] observed the advantage of core-rim structure in improving creep properties of ZrB2 by adding WSi2. The addition of Cr, Nb, Ti and Ta-borides were also attempted to improve oxidation resistance of ZrB2-SiC composites [16,21,22,23]. Some researchers have added transition metals in the form silicides to ZrB2 and observed the combined effect of transition metal and silicon on the oxidation behavior of the composite [16,22,24,25,26]. A comparative study revealed that MoSi2 is better reinforcement to improve oxidation resistance of ZrB2 than all the other transition metal silicides, as MoSi2 resulted more stable microstructural phases [26]. Feng et al. [27] studied the oxidation behavior of ZrB220SiC-5WC (vol%) composite at 1600 °C. It was reported that the combined addition of SiC and WC enhanced the oxidation resistance of ZrB2 composite than the ZrB2 solely added with either SiC or WC. The  ∗ Corresponding author. E-mail addresses: gbraju@nitw.ac.in, gbraju121@gmail.com (B.R. Golla).  https://doi.org/10.1016/j.ceramint.2019.04.076 Received 5 March 2019; Received in revised form 8 April 2019; Accepted 9 April 2019  Available online 10 April 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  Fig. 1. SEM and XRD patterns of starting powders (a and b) ZrB2, (c and d) SiC and (e and f) Ta, respectively.  Table 1  Relative density, grain size and lattice parameters of ZrB2 and (Zr, Ta)B2 solid solution with varying amount of  tantalum addition (RDRelative density).  Sample Id  Experimental density (g/cc)  RD from SEM images  (ZrB2) Core size (μm)  ZS-2.5Ta ZS-5Ta ZS-10Ta  5.56 5.65 5.87  99.11 99.11 99.63  2.76 ± 0.11 2.84 ± 0.08 2.65 ± 0.09  (Zr, Ta)B2 Rim size (μm)  4.17 ± 0.10 4.13 ± 0.14 3.61 ± 0.11  ZrB2  c (A°)  3.5314 3.5299 3.5304  (Zr, Ta)B2  a (A°)  c (A°)  a (A°)  3.1523 3.1655 3.1612  3.3386 3.4640  3.2075 3.1494  addition of transition metals or transition metal compounds to diborides modified its microstructure in the form of core-shell/rim structure and these characteristic features reportedly increased the high temperature mechanical properties as well [20,28,29]. In the literature, the transition metal silicides, carbides, nitrides or borides have been explored to improve properties of ZrB2-based composites [18-26,30]. Also, recently the effect transition metals on thermal properties of ZrB2 has been explored [31]. In this work, for the  first time we made an attempt to explore the effect of tantalum (Ta) transition metal addition on the microstructure and oxidation behavior of ZrB2-20vol% SiC. It has to be noted here that ZrB2-20vol% SiC has been selected as a base composition as it was proved to be optimal composite with better mechanical and oxidation properties. The ZrB220vol% SiC-Xwt.% Ta (X = 2,5, 5 and 10) composites were processed via Spark Plasma Sintering (SPS) (Temperature: 1900 °C, Pressure: 50 MPa; Sintering time: 3 min). The oxidation behavior of samples was  13800  \\x0c', \"S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  weight change, oxide layer crostructure analysis.  thickness, surface, and cross-sectional mi 2. Experimental procedure  The ZrB2 (purity > 98.2%, H.C. Starck Grade B GmbH and Co., Goslar Germany), SiC (purity > 99.8%, Alfa Aesar) and Ta (purity > 99.99%, India) powders were used to prepare the ceramic pellets. To process ZrB2-20vol% SiC-Xwt.% Ta (X = 2,5, 5 and 10) composites, appropriate amount of powders were weighed and mechanically mixed in a planetary ball mill (Fritsch Pulverisette 6, Germany) using silicon nitride media for 6 h in toluene at 200 rpm with balls to powder weight ratio at 2:1. Subsequently, the slurries were dried in a rotary evaporator. The dried powders were sintered using Spark Plasma Sintering (SPS) (Model: SPS 25-10, GT Advanced Technologies, USA) at a temperature of 1900 °C for 3 min under 50 MPa pressure in a vacuum (10−3 Pa). For convenience, the SPSed ZrB2-20vol% SiC-Xwt.% Ta (X = 2,5, 5 and 10) composites were designated as ZS-2.5Ta, ZS-5Ta, and ZS-10Ta. The detailed sintering schedule and more processing details can be found elsewhere [30]. The sintered specimens were polished and ultrasonically cleaned using acetone. The samples final densities were measured by the Archimedes method. The oxidation behavior of the ZrB2-based composites was evaluated using box type furnace (Nabertherm, MoSi2 heating element) exposing (2 × 4 × 8 mm3) the rectangular ZrB2 coupons in stagnant air at 1500 °C for 10 h. The coupons were cleaned in acetone using an ultrasonic bath and were dried and carefully measured the weight before and after the oxidation test. The cleaned coupons were inserted in an ultra-pure alumina crucible with the minimum contact area, and then crucible was placed at the center of the furnace. The microstructure before and after oxidation test was characterized using scanning electron microscopy (SEM, TESCAN VEGA 3 LMU) coupled to an energy dispersive spectroscopy (EDS, Oxford Instruments). The phases present in the powders, sintered and oxidized samples were characterized by Xray diffraction (XRD, PANalytical X'PertPro, Holland, CuKα = 1.5405A°). The microstructural features like grain size, the volume fraction of phases and oxide layer thickness etc. were evaluated using ImageJ software package (ImageJ 1.51j8, National Institute of Health, USA).  3. Results and discussion  3.1. Microstructure of starting powders  Fig. 1 shows the microstructure of ZrB2, SiC, and Ta initial powders that were used in the processing of ZrB2-SiC-Ta composites. The ZrB2 powders observed to have platelet shape with particle size varying in the range of 0.5-4.0 μm. The SiC powders are fine and in the agglomerated form with particle size ranging from 0.3 to 2.6 μm. On the other hand, the Ta powders are of mixed shape (rod-like, spherical and platelet) and its particle size vary between 0.6-7 μm. The corresponding XRD patterns of starting powders is also presented in Fig. 1 b, d, and f. It consists of only the respective crystalline phases and no other additional phases could be found within the detection limits of XRD.  3.2. Microstructure of sintered samples  Table 1 shows the densification and grain size of SPSed ZrB2-SiC-Ta composites. The sintered density of ZrB2-SiC-Ta composites increased from 5.56 to 5.87 g/cc, it was mainly due to the increasing addition of high density Ta to ZrB2-20vol% SiC (ZS). Irrespective of the amount of Ta addition, all the samples could be densified to more than 99% theoretical density (ρth) or relative density (RD). It shows that the selected SPS parameters are suitable sintering conditions for obtaining full density of ZS-Ta composites. The RD of samples were estimated by measuring the residual porosity from the SEM micrographs using  Fig. 2. Microstructures of multi-stage spark plasma sintered (a) ZS20-2.5Ta, (b) ZS20-5Ta and (c) ZS20-10Ta composites at 1900 °C, 50 MPa for 3 min.  studied by exposing the samples to a temperature of 1500 °C in the stagnant air for 10 h. The microstructural phase analysis of the samples before and after oxidation tests were investigated using XRD and SEMEDS. The oxidation resistance of samples was evaluated by measuring  13801  \\x0c\", 'S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  Fig. 3. Microstructure of (a) ZS-5Ta at high magnification showing the different phases, confirmed using EDS elemental analysis, which are (b) structure, (c) ZrB2 core and (d) (Zr, Ta)C and (e) SiC phases.  (Zr, Ta)B2 shell  Fig. 4. XRD patterns of Spark Plasma sintered (a) ZrB2-SiC-Ta composites (b) showing peak shifts, which confirms the solid solution formation (example: ZS-10Ta).  ImageJ software. The low magnification SEM images show hardly any traces of porosity and indicates its full density (Fig. 2). It also can be noticed that the dark contrasting phase (SiC) uniformly dispersed in the matrix phase. A closer look at the microstructure reveals core and rim structure of the matrix phase. As it was recorded in Table 1, the size of  (2.65-2.84 μm) and (Zr, Ta)B2 rim (3.61-4.17 μm) varied ZrB2 core narrowly. At the sintering temperature, Ta dissolves into ZrB2 matrix and forms the (Zr, Ta)B2 solid solution rim phase. As the amount of Ta increases, the amount of solid solution phase increases, which controls the growth of ZrB2 grains. The ZrB2 grain size after sintering remains  13802  \\x0c', \"S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  ZS with varying amounts of Ta show microstructural phases such as ZrB2, (Zr, Ta)B2 and SiC; and the boride matrix exists as “core-rim” structure, in which grains are comprised of ZrB2 core surrounded by an isostructural (Zr, Ta)B2 rim solid solution phase (see Figs. 2 and 3). The amount of Ta in the rim structure was in the limits of 0.5-1.5 at%, as observed by EDS. Fig. 3 shows high magnification image of ZS-5Ta sample with EDS elemental analysis. It depicts ZrB2, (Zr, Ta)B2, SiC and small fraction of bright (Zr, Ta)C phases, which was confirmed by EDS. The X-ray diffraction pattern of the as-sintered ZrB2-based composites is shown in Fig. 4. The XRD revealed the presence of ZrB2, (Zr, Ta)B2, SiC major crystalline phases along with minor amounts of ZrO2 and (Zr, Ta)C phases. In ZS-5Ta and ZS-10Ta samples, some peaks of monoclinic ZrO2 were observed which were formed due to pickup of oxygen during handling. In fact, similar observation of ZrO2 phase formation was reported by other researchers as well [32,33]. In case of ZS-5Ta and ZS10Ta samples, other phase was clearly observable at the right side of ZrB2 peaks, which was identified as hexagonal (Zr, Ta)B2 solid solution phase with reduced unit cell parameters. This phase was observed in ZS-2.5Ta sample at higher diffraction angle of 80°. The lattice parameters of ZrB2 and (Zr, Ta)B2 solid solution are reported in Table 1. The c/a ratio of the solid solution was smaller than ZrB2. It indicates the contraction of the unit cell, due to the smaller atomic radius of tantalum (0.143 nm) compared to zirconium (0.160 nm). So, according to the Bragg's law, the diffraction peaks of solid solution phase shifted towards higher angles (right side), as is shown in Fig. 4b. Monteverde [1] also reported change in the lattice parameters of core and rim structure of ZrB2-2.3MoSi2 and ZrB2-15SiC-2MoSi2 (vol%). The core-rim structure in those composites was observed due to the reaction between MoSi2 and surface oxides of ZrB2 [1]. The rim was composed of (Zr, Mo)B2 solid solution phase. The ZrB2 core and (Zr, Ta) B2 rim structure was also reported for ZrB2-15vol% TaSi2 and ZrB220SiC-(0-10)TaSi2 (vol%) composites [24,33]. In another work, Hu et al. [32] studied the microstructure and properties of spark plasma sintered ZrB2-(10-30)SiC-(10-20)TaSi2 (vol%). Interestingly, the corerim structure was observed for the ZrB2-based composites that were SPSed at 1600 °C and no such structure was present when the SPS was carried out at higher temperature of 1800 °C [32]. The formation of core-rim structure was also evidenced for WSi2, however, no such corerim structure was observed for ZrB2-15% ZrSi2 [26]. Only the presence of (Zr, TM)B2 solid solution phase was noticed when ZrB2 was added with small amount (1-1.2wt.%) of transition metals (TM) after hot pressing at 2150 °C, 32 MPa for 10 min [31]. It is mainly because of dissolution of the transition metals (Hf, Nb, W, Ti, Y) in to ZrB2. From this discussion it should be clear that the type of additive, its amount and processing conditions influence the microstructure of ZrB2-based composites. The fracture surfaces of SPSed ZrB2-SiC-Ta composites is presented in Fig. 5. It shows a combination of intergranular and transgranular fracture mode with good bonding between the grains without any hint of pores.  3.3. Oxidation of ZrB2-SiC-Ta composites at 1500 °C  The weight gain and oxide layer thickness of the samples with respect to the amount of tantalum addition after oxidation at a temperature of 1500 °C for 10 h are presented in Table 2. The oxidation properties of different ZrB2based composites reinforced with transition metal silicides, carbides or transition metals are also included for comparison purpose in the Table 2 [21,25,34,35,36,37,38]. The weight gain and oxide layer thickness were reduced with increasing amount of tantalum. The total oxide layer thickness of the composites was measured using cross-sectional SEM images of oxidized samples. From Table 2 it can be realized that the weight of Ta added ZrB2-20vol% SiC composites decreased from 22.91 to 18.77 mg/cm2 after isothermal oxidation at 1500 °C. Likewise, the oxide layer thickness (based on the SEM of the cross-sectional oxidation samples) of the ZrB2-SiC composites also decreased considerably from 401 to 195 μm with the addition  Fig. 5. SE images of 10Ta sample.  fracture surfaces of  (a) ZS-2.5Ta,  (b) ZS-5Ta and (c) ZS more or less similar to that of starting ZrB2 powders size (avg. particle ZrB2 size 2.5 μm). From Table 1, it also should be clear that the lattice spacing of the ZrB2 core and (Zr, Ta)B2 rim were varying and is due to the elastic and thermal mismatch during cooling from the sintering temperature.  13803  \\x0c\", 'S.K. Thimmappa and B.R. Golla  Table 2  Ceramics International 45 (2019) 13799-13808  Comparison of weight change and oxide layer thickness of ZrB2-SiC-Ta composites with other ZrB2-based materials after oxidation (p.w.  indicates present work).  Composition (vol%)  Conditions (Temperature in °C, time in h)  Weight gain (mg/ cm2)  Total oxide layer thickness (μm)  Oxide phases  ZrB2-20SiC-2.5wt.%Ta (ZS-2.5Ta) ZrB2-20SiC-5wt.%Ta (ZS-5Ta) ZrB2-20SiC-10wt.%Ta (ZS-10Ta) ZrB2-20SiC ZrB2-20SiC -5TaSi2 ZrB2 -20SiC -20TaSi2  ZrB2 ZrB2-10TaSi2 ZrB2-20TaSi2 ZrB2-30TaSi2 ZrB2-5.62B4C-27.91SiC-3.32TaSi2 (mol.%) ZrB2-5.62B4C-27.91SiC-13.29TaSi2 (mol.%) ZrB2-8Ta5Si3  ZrB2-31Ta5Si3  ZrB2-30vol% SiC2wt.% Y ZrB2-30vol% SiC7wt.% Y ZrB2-4mol.% W ZrB2-8mol.% W TiB2-2.5wt.% WSi2 TiB2-5wt.% WSi2 TiB2-10wt.% WSi2 TiB2-15wt.% WSi2  1500, 10 1500, 10 1500, 10 1627, 1.66 1627, 1.66 1627, 1.66  1400, 2 1500, 2 1500, 2 1500, 2 1500, 4 1500, 4 1400, 2  1400, 2  1700, 1 1700, 1 1600, 0.083 1600, 0.083 850, 64 850, 64 850, 64 850, 64  22.91 19.15 18.77 4.3 5.2 0.8  18.8 44 22 9 8.1 11.8 14.5  13.2  32.5 24.5 2.1 2.3 13 4.7 6.3 3.8  401 384 195 52.98  5.97  300 600 60 5.2 8 115   500 200 30 26  SiO2, ZrO2, TaZr2·75O8 SiO2, ZrO2, TaZr2·75O8 SiO2, ZrO2, TaZr2·75O8 ZrO2, SiO2 ZrO2, SiO2 ZrO2, SiO2, Ta(C, B) ZrO2 ZrO2, TaZr2·75O8 ZrO2, TaZr2·75O8 ZrO2, TaZr2·75O8 ZrO2, TaC ZrO2, TaC ZrO2, Ta2O5, SiO2, TaZr2·75O8 ZrO2, Ta2O5, SiO2, TaZr2·75O8 ZrO2, SiO2 ZrO2, SiO2, Y2Si2O7 ZrO2 ZrO2 TiO2, SiO2 TiO2, SiO2 TiO2, SiO2 TiO2, SiO2  Ref.  p.w. “ “  25 “ “  34 “ “ “  21 “  35  “  36 “  37 “  38 “ “ “  3.4. Microstructural characterization  Fig. 6 shows the XRD of the ZrB2 sample surfaces after oxidation at 1500 °C for 10 h. The presence of monoclinic ZrO2 and orthorhombic TaZr2.75O8 crystalline phases on the ZrB2 samples. It is evident that TaZr2.75O8 phase in ZrB2 samples increased with more amount of Ta addition. Peak broadening at about 22o indicate presence of amorphous SiO2. The SEM oxidized ZrB2 samples can be seen in Fig. 7. Spherical and dendritic ZrO2, dark SiO2 matrix phase and coarse spherical TaZr2.75O8 were noticeable on the oxidized ZrB2-based composites. It is also observed that as the amount of Ta increases, ZrO2 grains transformed from spherical to dendritic shape. SEM-EDS of ZS-5Ta sample after the oxidation is shown in Fig. 8. The presence of ZrO2, SiO2 and TaZr2.75O8 phases were confirmed by EDS. The formation of TaZr2.75O8 phase can be attributed to oxidation reactions (1) and (2). The thermodynamic feasibility of these reactions were possible at the temperature of 1500 °C [26,39,40].  (Ta,Zr)B2(s) + O2(g) → TaZr2.75O8(s)  (Ta, Zr)C(s) + O2(g) → TaZr2.75O8(s)  (1)  (2)  Further the cross-sectional SEM of ZrB2-SiC-Ta composites was carried out to understand the oxide layer presence in the samples (Fig. 9). It is evident that all the ZrB2 samples consists of three different layers: thick and dense outer SiO2 layer, intermediate SiC depleted layer and unreacted bulk. The thickness of oxide layer decreased considerably with higher amount of Ta addition to ZrB2-20vol% SiC. The SEM-EDS of cross-sectional SEM-EDS of ZS-10Ta sample after oxidation is presented in Fig. 10. The EDS analysis from all the three layers clearly confirms the presence of SiO2, SiC depleted layer and unreacted bulk in the oxidized samples. It is interesting to note the presence of SiC depleted layer even though there is no indication of any defects presence at or near the external SiO2 layer (see Fig. 10).  3.5. Oxidation mechanism  During oxidation, initially ZrO2 and B2O3 forms due to the oxidation of ZrB2 in the temperature range of 800-1200 °C. In particular, B2O3 is protective at low temperatures and the evaporation of B2O3 starts at  Fig. 6. XRD patterns of tantalum (2.5, 5 and 10wt.%) added ZrB2-20vol% SiC (ZS) composites after oxidation at 1500 °C for 10 h.  of Ta. It indicates the usefulness of Ta in improving the oxidation resistance of ZS. The reduction in the weight of oxidized samples with increasing amount of tantalum is due to either evaporation of gaseous phases at the oxidation temperature or less diffusion of oxygen into the surface of the sample. The Ta addition to ZrB2-SiC composite may help to develop immiscibility in the top passive layer, which has a higher melting point and retards the diffusion of oxygen and thereby no chance of evaporation of silica at 1500 °C. The oxide layer thickness also decreases with increasing amount of tantalum due to passive behavior of the top silica layer, which retards the diffusion of oxygen further into the surface. It also observed that as the amount of tantalum increases the thickness of the middle layer decreases due to the protective nature of the top silica layer.  13804  \\x0c', 'S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  oxidized to form SiO2, which protects the composite from further oxidation. The SiO2 scale formed on ZrB2-SiC composite, which is stable at high temperatures than the boria layer formed because of low volatility of liquid silica glass compared to boria at these conditions. The overall oxidation of ZrB2-SiC composites takes place by the oxidation reactions (1) to (7).  ZrB2(s)+5/2O2(g) → ZrO2(s)+B2O3(l)  B2O3(l) → B2O3(g)  SiC(s) + 3/2O2(g) → SiO2(l) + CO(g)  SiC(s) + O2(g) → SiO(g) + CO(g)  SiC(s)+ 2SiO2(l) → 3SiO(g) + CO(g)  (3)  (4)  (5)  (6)  (7)  During prolonged holding (10 h) at the oxidation temperature of 1500 °C, SiO(g) forms by the oxidation (via reaction 6 and 7) and thus it lead to the formation of SiC depleted layer. The observation of SiC depleted layer has been reported by several researchers for ZrB2-SiC composites [5]. Interestingly, the presence of SiC depleted layer was evidenced for graphite, TaSi2, TaB, TiB2 reinforced ZrB2-SiC composites [4,5,12,16,22,24,25]. However, the La2O3, LaB6, WC and Si3N4 added ZrB2-SiC composites did not show any indication of SiC depleted layer [7,22,27,29,30]. Depending on the oxidation conditions the SiC depletion was reported for ZrB2-SiC-LaB6 conditions [22,23]. As mentioned above, the SiO2 layer thickness decreased with increasing the amount of tantalum to ZrB2-SiC. The possible reason for decreasing the SiO2 layer thickness due to glass immiscibility developed by Ta cations in liquid silica. Silvestroni and Kleebe [24] proposed the Ta presence improves efficacy of oxidation resistance of ZrB2-15TaSi2 due to increased cation field strength by promoting glass immiscibility and retards the diffusion of oxygen. In case of ZrB2-15vol% TaSi2 composites, it was reported that ZrO2, TaZr2·75O8 and TaB2 phases present after oxidation at 1500 °C and ZrO2, TaZr2·75O8 and Ta2O5 phases at oxidation temperature of 1650 °C for 15 min in air [24]. The formation of Ta2O5 phase inside ZrO2 grains develops the cracks and spallation of the oxide layer from the inner bulk, which reduces the oxidation resistance at above 1650 °C. Nevertheless, TaB2 and Ta2O5 phases were not detected in this current work, instead of that TaZr2·75O8 was observed due to continuous oxidation for 10 h. In a different work, presence of ZrO2 and TaZr2·75O8 phases were observed in oxidized ZrC-TaSi2 composites [39,40]. However, Levine et al. [16] observed only ZrO2 and SiO2 phases in oxidized (at 1627 °C for 10 min) ZrB2-20vol% TaSi2 composite without any traces of tantalum containing phases. He et al. [36] studied the effect of Yttrium (Y) addition on the oxidation behavior of ZrB2-SiC composite. The Y dissolved in the silica layer, thereby it improved oxidation resistance of the ZrB2-based composites by forming Y2Si2O7 phase. Dehdashti et al. [41] studied the effect of Mo, W and Nb addition on the oxidation behavior of monolithic ZrB2 ceramic. The thickness of oxide layer is above 300 μm for Mo, W and Nb reinforced ZrB2 after oxidation at 1600 °C for 3 h. The oxidation properties of different ZrB2-based composites are compared in the Table 2 [21,25,34,35,36,37,38]. The weight thickness (52.98-5.97 μm) of TaSi2 (4.3-0.8 mg/cm2) and oxide layer reinforced ZrB2-20SiC decreased after oxidation at 1627 °C for 1.6 h [25]. The improvement in the oxidation resistance was attributed to glass immiscibility and high viscosity of glass with Ta. Similarly, the addition of TaSi2 also considerably lowered the weight (44-9 mg/cm2) and oxide layer thickness (300-60 μm) of ZrB2 after oxidation at 1400/ 1500 °C for 2 h [34]. The addition of Yttrium to ZrB2-30SiC also re(32.5-24.5 mg/cm2) duced its weight and oxide layer thickness (500-200 μm) after oxidation at 1700 °C for 1 h [36]. A comparison of present work results with the literature data indicates the usefulness of Ta in improving oxidation resistance of ZrB2-20SiC. Since the oxidation test conditions and materials compositions are different, the underlying  Fig. 7. Surface morphology of (a) ZS-2.5Ta, (b) ZS-5Ta and (c) ZS-10Ta samples after oxidation at 1500 °C for 10 h.  temperatures higher than 1100 °C and left porous ZrO2. When temperature reaches to 1500 °C, the outer oxide scale composition changes significantly. Due to high vapor pressure, B2O3 evaporates and SiC  13805  \\x0c', 'S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  Fig. 8. Surface morphology of (a) ZS-5Ta sample and its corresponding phases EDS i.e., (b) ZrO2, (c) TaZr2·75O8 and (d) SiO2 after oxidation at 1500 °C for 10 h.  Fig. 9. Cross-sectional SEM of (a) ZS-2.5Ta, (b) ZS-5Ta and (c) ZS-10Ta samples after oxidation at 1500 °C for 10 h.  oxidation mechanisms will be different.  4. Conclusions  a. More than 99% theoretical density was achieved for ZrB2-20vol% SiC-Xwt.% Ta (X = 2,5, 5 and 10) composites after spark plasma sintering at 1900 °C, 50 MPa for 3 min. b. The microstructure of sintered samples consists of ZrB2 core, (Zr, Ta) B2 rim, SiC as major crystalline phases and minor amounts of ZrO2 and (Zr, Ta)C phases. c. The ZrB2-SiC-Ta composites composed of crystalline ZrO2, TaZr2·75O8 and amorphous SiO2 phases after isothermal oxidation at  1500 °C for 10 h. d. The cross-sectional SEM-EDS of oxidized samples revealed threelayered architecture: the top passive SiO2 layer, intermediate SiC depleted layer and bottom unreacted bulk. e. The decrease in weight and oxide layer thickness of ZrB2-20vol% SiC with increasing amounts of tantalum confirm its improved oxidation resistance.  Acknowledgement  The financial support by Science and Engineering Research Board of Department of Science and Technology, Government of India is  13806  \\x0c', 'S.K. Thimmappa and B.R. Golla  Ceramics International 45 (2019) 13799-13808  Fig. 10. (a) Cross-sectional SEM-EDS of ZS-10Ta sample after oxidation at and 1500 °C for 10 h and (b), (c) and (d) showing EDS spectra of top (SiO2), middle (SiCdepleted layer) and unreacted bulk.  gratefully acknowledged.  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2019.04.076 .  References  [2]  [1]  [3]  F. Monteverde, The addition of SiC particles into a MoSi2-doped ZrB2 matrix: effects on densification, microstructure and thermo-physical properties, Mater. Chem. Phys. 113 (2-3) (2009) 626-633. J. Zou, S.K. Sun, G.J. Zhang, Y.M. Kan, P.L. Wang, T. Ohji, Chemical reactions, anisotropic grain growth and sintering mechanisms of self-reinforced ZrB2-SiC doped with WC, J. Am. Ceram. Soc. 94 (5) (2011) 1575-1583. S.Q. Guo, Densification of ZrB2-based composites and their mechanical and physical properties: a review, J. Eur. Ceram. Soc. 29 (6) (2009) 995-1011. [4] A. Rezaie, W.G. Fahrenholtz, G.E. 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Majumdar, B. Basu, Phase stability, hardness and oxidation behaviour of spark plasma sintered ZrB2-SiC-Si3N4 composites, Ceram. Int. 45 (7 PART A) (2019) 9061-9073. [31] D.L. McClane, W.G. Fahrenholtz, G.E. Hilmas, Thermal properties of (Zr,TM)B2 solid solutions with TM = Hf, Nb, W, Ti, and y, J. Am. Ceram. Soc. 97 (5) (2014) 1552-1558. C. Hu, Y. Sakka, H. Tanaka, T. Nishimura, S. Guo, S. Grasso, Microstructure and properties of ZrB2-SiC composites prepared by spark plasma sintering using TaSi2 as sintering additive, J. Eur. Ceram. Soc. 30 (12) (2010) 2625-2631. Y. Yang, Y. hai Qian, J. jun Xu, M. shuan Li, Effects of TaSi2 addition on room temperature mechanical properties of ZrB2-20SiC composites, Ceram. Int. 44 (14) (2018) 16150-16156. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, High-temperature chemistry and oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3, and TaSi2, J. Am. Ceram. Soc. 91 (7)  [32]  [33]  [34]  [38]  [36]  [35]  (2008) 2250-2257. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, A.H. Smith, Properties of ceramics in the system ZrB2-Ta5Si3, J. Mater. Res. 21 (10) (2006) 2593-2599. J. He, Y. Wang, L. Luo, L. An, Oxidation behaviour of ZrB2-SiC (Al/Y) ceramics at 1700 °C, J. Eur. Ceram. Soc. 36 (15) (2016) 3769-3774. [37] M.K. Dehdashti, W.G. Fahrenholtz, G.E. Hilmas, Effects of temperature and the incorporation of W on the oxidation of ZrB2 ceramics, Corros. Sci. 80 (2014) 221-228. T.S.R.C. Murthy, J.K. Sonber, C. Subramanian, R.C. Hubli, A.K. Suri, Densification, characterization and oxidation studies of TiB2-WSi2 composite, Int. J. Refract. Metals Hard Mater. 33 (2012) 10-21. L. Silvestroni, D. Sciti, M.B. Pichelin, L. Charpentier, Zirconium carbide doped with tantalum silicide: microstructure, mechanical properties and high temperature oxidation, Mater. Chem. Phys. 143 (1) (2013) 407-415. L. Charpentier, M.B. Pichelin, D. Sciti, L. Silvestroni, High temperature oxidation of Zrand Hf-carbides: influence of matrix and sintering additive, J. Eur. Ceram. Soc. 33 (15-16) (2013) 2867-2878. [41] M.K. Dehdashti, W.G. Fahrenholtz, G.E. Hilmas, Effects of transition metals on the oxidation behavior of ZrB2 ceramics, Corros. Sci. 91 (2015) 224-231.  [39]  [40]  13808  \\x0c']"
},{
  "_id": 56,
  "PDF": "Effect of the ZrB2 content on the oxygen blocking ability of ZrB2-SiC coating at 1973K.pdf",
  "Text": "['Contents lists available at ScienceDirect   Journal of the European Ceramic Society   journal homepage: www.elsevier.com/locate/jeurceramsoc   Original Article   Effect of the ZrB2 content on the oxygen blocking ability of ZrB2-SiC coating  at 1973K   Xuanru Ren *, Hongao Chu, Keyou Wu, Anni Zhang, Menglin Huang, Can Ma, Haifeng Liu,  Peizhong Feng *   School of Materials Science and Physics, China University of Mining and Technology, Xuzhou 221116, China     A R T I C L E I N F O   A B S T R A C T   Keywords:  ZrB2  Coating  SPS  Oxidation inhibition  Nanocrystal dispersion   ZrB2-SiC coatings were prepared by spark plasma sintering method. The effect of ZrB2 content ranging from  20wt.% to 80 wt.% on the oxidation behavior of the ZrB2-SiC coatings was investigated at 1973K in air. In the  initial oxidation stage (<30 min), although the increased ZrB2 content with high oxidation activity aggravated  the oxidation of coating, the slackened SiC consumption enhanced the oxygen blocking ability of the coating  itself. With increasing oxidation time, owing to sufficient dispersion of Zr-oxides nanocrystal particles in the self generated glass, the glassy layer on the 60ZrB2-40SiC coating exhibited the best oxygen blocking ability. The  dispersed Zr-oxides nanocrystal particles conduced to broaden crystalline areas and increased the viscosity of the  glassy layer, restraining the oxygen permeability and weakening the damage of oxygen to the glassy layer. With  increasing oxidation time, the dispersion degree of Zr-oxides nanocrystal particles would strengthen the  oxidation protection efficiency.     1.  Introduction   Carbon structural materials [1-3] are widely used in aerospace industry as thermal structure materials, owing to their low density, high  temperature mechanical properties, excellent high temperature and  thermal shock resistance [4-7]. However, carbon structural materials  are prone to react with oxygen at elevated temperature, causing fatal  damaging and erosion [8-10]. Hence, the oxidation resistance of the  carbon structural materials becomes critical when reusability becomes  an issue [11-14]. To inhibit the oxidation of carbon structural materials,  ZrB2-SiC ceramics have been considered owing to the high melting  (>3000℃) and excellent ablation resistance of the boride, and the  protective action of the self-generated Zr-B-Si-O glassy layer [15-22].  However, although the ZrB2-SiC coatings exhibit great oxidation  protective ability below 1800 K, when the service temperature is higher  than 1973 K, the oxidation mode of SiC will pass from passive oxidation  to active oxidation [23]. The aggravated oxidation consumption of SiC  will not only destroy the oxygen barrier structure of the coating itself,  but also generate numerous gaseous by-products and damage the  integrity of the glassy and subscale  layer, thus  increasing oxygen  diffusion channels in the coating. Hence, the dramatically raised application temperature puts forward harsh requirements for the oxidation   resistance of the ZrB2-SiC coating.  Since the coating itself is the only barrier to oxidation for the carbon  substrate, it is necessary to prepare ZrB2-SiC coatings with dense and  crack-free structure to prevent the inward penetration of oxygen, which  is closely related to the preparation process of the coating. In view of the  high melting point of ZrB2 and the poor sintering characteristics of SiC,  pressure-less sintering techniques, such as pack cementation [24],  in-situ reaction method [25], liquid phase sintering [26], have attracted  much attention to synthesize the ZrB2-SiC coatings. However, the  pressure-less sintering techniques usually need to reach temperatures  above 2000 K to achieve densification, the exorbitant sintering temperature is prone to cause excessive grain growth and energy consumption. Hence,  the hot-pressing  technology  conducted at  low  temperature with rapid sintering speed becomes very attractive for  production of simple shapes. Spark plasma sintering (SPS) method  [27-30] is even a more promising densification technique. Recently, we  have used the SPS method to prepare MoSi2-SiC coatings [31] and  MoSi2-MoB-ZrO2 coatings [32], presenting great potential. Nevertheless, no work has been reported about the preparation of ZrB2-SiC  coatings by SPS method.  Both static and dynamic oxidation are important parameters for the  evaluation of a coating performance. The capability to effectively hinder   * Corresponding author.  E-mail addresses: XuanruRen@163.com (X. Ren), fengroad@163.com (P. Feng).    https://doi.org/10.1016/j.jeurceramsoc.2020.10.036  Received 12 June 2020; Received in revised form 16 October 2020; Accepted 19 October 2020         \\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        oxygen penetration depends on the oxidation consumption degree of the  coating itself and on the oxidation resistance of the self-generated glassy  layer, which is closely related to the coating components. In our previous work, we have found that owing to the dispersion of Zr-oxides in Zr B-Si-O glassy layer, increasing ZrB2 content in the coating inhibits the  oxygen penetration in a broad temperature region from 298 K to 1773 K  [33]. To note that, due to the good compatibility with carbon matrix,  and the self-generated glass, SiC is necessary for the oxidation resistance  of the ZrB2-SiC coating. However, its content must be carefully designed  in order to enable the formation of a protective glass, but avoid excessive  oxide embrittlement once the conditions of temperature and oxygen  pressure boost SiC active oxidation.  The oxidation of the pure ZrB2-SiC coatings and the carbon substrate  show opposite trend of mass change, and therefore, the oxidation curves  of the coating samples are not authentic enough to reflect the real oxygen blocking effect of the ZrB2-SiC coatings. Hence, it is of great significance to separate and distinguish the intrinsic oxidation behavior of  the coatings. To date, seldom systematic work has been reported on the  oxidation behavior of ZrB2-SiC coating above 1973 K.  Thus, in this work, ZrB2-SiC coatings with different ZrB2 contents  ranging from 20 wt.% to 80 wt.% were prepared by the SPS method. To  study the authentic oxygen blocking ability of the ZrB2-SiC coatings, a  bulk reference-coating sample-substrate sample data collection mode  was set up, in which the ZrB2-SiC ceramic bulk reference samples and   Fig. 1. XRD patterns of  SPS method.   the ZrB2-SiC multiphase   coating prepared by   Fig. 2. Surface SEM images of the ZrB2-SiC coatings with different ZrB2 contents: (a) 20 wt.%; (b) 40 wt.%; (c) 60 wt.%; (d) 80 wt.%; (e) high-resolution TEM images  of (d); (f) density of (a)-(d).   JournaloftheEuropeanCeramicSociety41(2021)1059-10701060\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 3. Polished cross-section SEM images of the ZrB2-SiC multiphase coatings with different ZrB2 contents: (a) 20 wt.%; (b) 40 wt.%; (e) 60 wt.%; (d) 80 wt.%; (e)-  (h) magnification of (a)-(d).   with alcohol and dried at 353 K for 2 h. In order to explore the influence  of different ZrB2 contents on the oxidation resistance of the coatings,  four groups of experiments were carried out with ZrB2 powder (Beijing  Licheng Innovation Metal Materials Technology Co., Ltd, <200 μm,  ≥99.9 %) and SiC powder (Beijing Licheng Innovation Metal Materials  Technology Co., Ltd, 200 meshes, ≥99.9 %) as raw materials. The  coating formulations were 20 wt.% ZrB2-80 wt.% SiC, 40 wt.% ZrB2 60 wt.% SiC, 60 wt.% ZrB2-40 wt.% SiC and 80 wt.% ZrB2-20 wt.% SiC.  The oxidation resistant coatings were prepared by spark plasma  sintering (lab-125, sinter land Inc., Japan) under vacuum atmosphere.  The raw material powders were weighed and put in agate tank first, then  ball milled with agate balls at the speed of 200 r/min for 5 h to make the  raw coating materials uniformly mixed. Next, 1.9 g mixed powders were  put onto the graphite mould with a pressure head, making the treated  graphite matrix completely covered with the mixed powders. The  components were sintered rapidly at 1773 K with external pressure of  30 MPa. The heating rate was 200 K/min and the holding time 5 min. In  order to investigate the oxidation resistance of different ZrB2 contents,  the pure ceramic bulk reference samples without carbon substrate were  sintered under the same conditions for comparative analysis. The  graphite paper around the coatings was grinded off after sintering to  eliminate its influence on the oxidation resistance of the coating.   2.2.  Isothermal oxidation test at 1973 K   The ZrB2-SiC coating samples and the corresponding bulk reference  samples with different composition were introduced into the resistance  furnace (SGM M4/18AE, Luoyang SIGMA high temperature electric  furnace Co., Ltd.) and isothermal oxidized at 1973 K for 100 min. The  samples were taken out to weigh every 10 min during the oxidation  process up to 100 min, the weight changes were accurately measured  with a balance. The oxidation of each composites was repeated twice to  have reliable results to build the curves of weight change and weight  change rate, according to Eqs. (1) and (2). Based on the recorded data  changes, the oxidation protective ability of the coatings was further   Fig. 4. XRD patterns of ZrB2-SiC coatings oxidized at 1973 K in air for 100 min.    the corresponding coating samples were simultaneously prepared by the  SPS method for comparison. The effect of ZrB2 content on the oxygen  blocking ability of the ZrB2-SiC coatings and protective action of the  different Zr-B-Si-O glassy layers at 1973 K are here investigated.   2.  Experimental procedure   2.1.  Specimen preparation   The dimensions of the graphite matrix (Qingdao Baofeng graphite  products Co., Ltd., ≥99.995 %) prepared by mould pressing technology  was Ф12 mm × 2.5 mm with density of 1.6 g/cm3. Before applying the  coatings, the graphite matrices were polished, ultrasonically cleaned   JournaloftheEuropeanCeramicSociety41(2021)1059-10701061\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 5. Weight change curves of ZrB2-SiC coatings (a) and bulk reference samples (b) oxidized at 1973 K; (c) and (d) the corresponding weight change rates of (a) and  (b); Aactive oxidation stage; B-steady oxidation stage.   discussed.   mk,j  G =  mk,0  (cid:0)  mk,0  3. Result and discussion   (1)    3.1. Microstructure analysis before oxidation   Where G is the weight gain percentage; mk,j is the weight per unit area of  k formulation sample at j time; mk,0  is the initial weight of the k  formulation samples.   mk,j  (cid:0)  Vk =  mk,j  (cid:0)  1  tj  (2)    Where Vk is the rate of weight change; mk,j is the weight per unit area of k  formulation sample at  j time;  is the time  interval of weighing  (tj = 10 min).   tj   2.3.  Samples characterization   X-ray diffraction (XRD, Bruker AXS, Germany) was employed to  analyze the phase composition of the ZrB2-SiC coatings before and after  isothermal oxidation. The surface and cross-section morphology of the  sample were analyzed by scanning electron microscopy (SEM, JSM6460,  JEOL, Japan) equipped with the energy dispersive spectroscopy (EDS).  Transmission electron microscopy (TEM, Tecnai G2 F20, FEI, America)  was used to analyze the micro-morphology and crystal structure of the  ZrB2-SiC coating fragment before oxidation and glassy layer fragment  after isothermal oxidation tests. To make the TEM specimens, first, some  pieces were scraped off the surface of the ZrB2-SiC coating samples and  the glassy layer samples; then, the scraps were grinded in a mortar for  half an hour using alcohol as the abrasive; next, the TEM samples can be  obtained by taking the supernatant of alcohol solutions with fragments.   Fig. 1 shows the XRD patterns of the ZrB2-SiC coatings where only  ZrB2 (JCPDS Card reference code #02-75-0964) and SiC phases (JCPDS  Card reference code #05-89-2214) are present with sharp shape and  high intensity, indicating the high crystallinity of the coatings. Moreover, with increasing the ZrB2 content, the XRD patterns presented  increased intensity of the ZrB2 peaks and decreased intensity of the SiC  peaks. However, a few peaks of C could be detected on the coatings due  to the residue of the graphite paper.  Fig. 2 (a)-(d) shows the SEM micrographs of the polished surfaces of  the ZrB2-SiC coatings, in which two kinds of phases could be observed.  By EDS analyses, the black and white phase were SiC and ZrB2,  respectively. With increasing the ZrB2 content, the coatings presented a  significant component change with enhanced compactness, and consistent with the XRD results. Moreover, no obvious crack was found in the  micro-structure. The high-resolution TEM image of ZrB2-SiC coatings  before oxidation is shown in Fig. 2 (e). The ZrB2 and SiC grains were well  connected, indicating good densification of the ZrB2-SiC coatings by SPS  method. Fig. 2 (f) shows the measured and theoretical density of the  ZrB2-SiC coating samples, the difference of which got smaller with  increasing the ZrB2 content, indicating that the density of the ZrB2-SiC  coating improved with increasing the ZrB2 content.  SEM micrographs of the polished cross-section of the ZrB2-SiC  coatings are shown in Fig. 3. The coating thickness ranged from 700 μm  to 900 μm. In addition, no obvious microcrack or apparent gap between  the coatings and the substrate existed in the cross-section of the coatings.  Moreover, as shown in Fig. 3 (e) (h), the cross-section images presented  similar microstructures, reflecting the uniform distribution of each  component in the coating.   JournaloftheEuropeanCeramicSociety41(2021)1059-10701062\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 6. Weight gain of the coatings (a) and bulk reference samples (b) in active oxidation stages; weight gain of the coatings (c) and bulk reference samples (d) in  steady oxidation stages; K values of the samples in the active (e) and steady (f) oxidation stage.   JournaloftheEuropeanCeramicSociety41(2021)1059-10701063\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 7. Oxidation protection efficiency curves of samples during oxidation  at 1973 K.   Fig. 8. XRD patterns of the 80 wt.% ZrB2-20 wt.% SiC coating samples after  oxidation at 1973 K for different times.   Fig. 9. Surface backscattered SEM images of the 80ZrB2-20SiC coating samples oxidized at 1973 K for different times: (a) 20 min; (b) 40 min; (c) 60 min; (d) 80 min;  (e) 120 min; (f)-(j) magnification of (a)-(e); (k)-(o) cross-section SEM image of (a)-(e); (p)-(t) corresponding EDS Mapping micrographs of O corresponding to  (a)-(e).   JournaloftheEuropeanCeramicSociety41(2021)1059-10701064\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 10. Surface EDS mapping of the 20 wt.% ZrB2-80 wt.% SiC coating after oxidation at 1973 K for 100 min.    3.2. Microstructure evolution of the coatings upon oxidation at 1973 K   Fig. 4 shows the XRD spectra of the ZrB2-SiC coatings oxidized at  1973 K in air for 100 min. m-ZrO2 (JCPDS Card reference code # 00 013-0307), t-ZrO2 (JCPDS Card reference code # 00-002-0733), and  ZrSiO4 phases (JCPDS Card reference code #02-83-1375) could be  observed after oxidation of the coatings. However, no obvious diffraction peak of SiO2 could be observed, reflecting the amorphous nature of  the SiO2 glass. The involved reaction equations are presented in Eqs.  (3)-(10) [34]. With increasing ZrB2 content, the intensity of the ZrO2  and ZrSiO4 increased. Furthermore, the intensity of SiC kept increasing  with  increasing ZrB2 content, suggesting  that  the  increased ZrB2  inhibited SiC active oxidation. Therefore, the composition changes in  the pristine coating and hence in the developed oxide layers would  greatly affect the oxidation protection ability of the coatings.    2ZrB2(s)+5O2(g) = 2ZrO2(s)+2B2O3(s)                                               B2O3(s)→B2O3(l)→B2O3(g)                                                                SiC(s)+2O2(g)=SiO2(s)+CO2(g)                                                         2SiC(s)+3O2(g) = 2SiO2(s)+2CO(g)                                                    ZrO2(s)+SiO2(s)=ZrSiO4(s)                                                                2SiO2(s)+SiC(s) = 3SiO(g)+CO(g)                                                     C(s)+O2(g)=CO2(g)                                                                          (3)    (4)    (5)    (6)    (7)    (8)    (9)    2C(s)+O2(g) = 2CO(g)                                                                    (10)   As shown in Fig. 5 (a) and (b), the weight gain curves of ZrB2-SiC  coating samples and bulk reference samples exhibited the same trend,  which gradually slowed down with extending the oxidation time over  30(cid:0) 40 min. Fig. 5 (c)-(d) shows the weight change rates of the samples.  The oxidation rates of all samples increased first and then decreased, and  finally reached a relatively stable state. According to the weight change  rate curves, the oxidation process of the ZrB2-SiC samples could be  roughly divided into two stages, the active oxidation stage A (<30 min)  and steady oxidation stage B (>30 min). However, no matter if the  curves of weight gain or weight change rate were considered, the values  relative to coating samples were smaller than those of the corresponding  bulk reference samples, owing to the graphite substrate consumption in  the remained exposed surfaces.  Considering the nearly linear oxidation behaviors of the samples, the  curves of all samples separated by oxidation active stage (up to 30 min)  and steady stage (40(cid:0) 100 min) in different oxidation stages were fitted  according to the first-order kinetics of oxidation equation (Δw = Kt + B),  as shown in Fig. 6 (a)-(d). In the equation, Δw (%) represents the weight  increase percentage of the samples, t (min) represents the cumulative  oxidation time. K represents the oxidation activation capacity coefficient of the samples, and the lower the K value is, the better the  oxidation resistance of the coating is. B is constants. In the active  oxidation stage, the main barrier to oxygen diffusion of the samples was  the coating itself. As shown in Fig. 6 (e), with raising the ZrB2 content,   JournaloftheEuropeanCeramicSociety41(2021)1059-10701065\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 11. (a) TEM images of the Zr-B-Si-O glassy layer fragments; (b) magnification of part A in (a); (c) high resolution TEM images of the nanocrystalline region; (d)  sketch of the Zr-oxides dispersion.   the K value of samples gradually increased, demonstrating that an increase in ZrB2 with higher oxidation activity would aggravate the  oxidation of the coating. However, as shown in Fig. 6 (f), in view of the  high oxidation activity of ZrB2, the K value of the samples in the steady  oxidation stage did not  increase as raising the ZrB2 content, but  decreased gradually. Accordingly, in the steady oxidation stage, the K  values of the coating and bulk samples in the steady oxidation stage  were significantly lower than that in the active oxidation stage, indicating that the formation of oxidation barrier layer was responsible for  the characteristic of the steady oxidation stage. Furthermore, the  inhibited oxidation consumption of the coating indicated that raising the  ZrB2 content conduced to enhance the stability and oxidation inhibition  ability of the glassy layer.  Considering the weight change differences between the coating and  the corresponding bulk reference samples, the protective ability of the  coatings was evaluated by considering the cumulative oxidation protection efficiency (ηk,j) of the coating and bulk samples based on Eq.  (11), whose curves are shown in Fig.7. The higher the ηk,j value is, the  better oxidation protective effect of the coating is. The ηk,j curves of all  samples showed a rapidly decreasing trend up to 30 min exposure and  then flattening, reflecting the different oxidation protection ability of  the coatings mostly at the beginning of the oxidation process. When the  oxidation time is 10 min, the ηk,j value of the coating containing 20 wt.%  ZrB2 was the lowest, indicating the violent oxidation of SiC at 1973 K  destroyed the oxidation protective effect of the coating. Moreover, in the  active oxidation stage, when the content of ZrB2 was above 40 wt.%,  increasing the ZrB2 content did not enhance the ηk,j value but decreased,   further testifying the higher oxidation activity of ZrB2 would weaken the  oxidation protective effect of the coating in the initial oxidation stage.  However, in the steady oxidation stage, the ηk,j curves presented opposite trend. For the coatings whose content of ZrB2 was below 40 wt.%,  the ηk,j value tended to be stable as increasing exposure time; while with  further raising ZrB2 content, the ηk,j value of the coatings tended to  present rising trend, demonstrating that raising ZrB2 content improved  the oxidation resistance of the ZrB2-SiC coatings with increasing serving  time.   {  ηk,j =  1 (cid:0)  (ms,k,j  (cid:0)  ms,k,0 ) (cid:0) (cid:0)  mp,j  (cid:0)  (mc,k,j  )  mp,0  }  (cid:0)  mc,k,0 )  × 100%  (11)    Where ηk,j is the oxidation protection efficiency of k formulation coating  sample at j time; mc,k,j and ms,k,j are the weight per unit area of k  formulation coating sample (c) and bulk reference sample (s) at j time;  mp,j is the weight per unit area of pure graphite substrate sample (p) at j  time. The subscript s, c and p indicate bulk sample, coating sample and  graphite substrate, respectively.   3.3. Oxidation behaviors of the coatings at 1973 K   To disclose the phase evolution of the oxidation products of the ZrB2 SiC coatings during oxidation, the XRD spectra of the 80 wt.%ZrB2 20 wt.% SiC coating samples oxidized at 1973 K in air for different times  are shown in Fig. 8. With increasing oxidation time, the diffraction peak  intensity of ZrO2 phases showed a trend of first strengthening and then   JournaloftheEuropeanCeramicSociety41(2021)1059-10701066\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Fig. 12. The macro-morphology of the ZrB2-SiC coating samples with different ZrB2 content after oxidation at 1973 K for 100 min: (a) 20 wt.%; (b) 40 wt.%; (c)  60 wt.%; (d) 80 wt.%; (e)-(h) magnification of (a)-(d); (i)-(l) cross-section SEM image of (a)-(d); (m)-(p) corresponding EDS mapping micrographs of O corresponding to (a)-(d).   dispersed in the glassy layer, which is caused by the rapid spread of the  generated fluid SiO2 glass resulting in a continuous glassy scale coating  shown in Fig. 2. With prolonging oxidation time, more and more white  Zr-oxide precipitated in the glassy layer, as shown in Fig. 9 (b)-(e) and  (g)-(j). As the oxidation time increased, the content of the Zr-oxides  particles increased as well, possibly caused by the transition from  ZrO2 particles to stable ZrSiO4 phase. Moreover, with extending oxidation time, the diameter of pores got smaller, indicating the dispersion of  Zr-oxides intensified the ability of the glass to inhibit the escape of gas  by-products. The corresponding cross-section EDS elemental mapping in  the coatings after oxidation are shown in Fig. 9 (k) (o). With increasing  the oxidation time, the thickness of the generated glassy layer resulted  120 μm, 148 μm, 168 μm, 180 μm and 185 μm, respectively. Hence, the  thickness of the glassy layer did not constantly increase, but tended to  achieve a steady state after a certain period. Accordingly, with extending oxidation time, the dispersion of Zr-oxides enhanced the stability of  the glassy layer to prevent oxygen permeation.  In order to further analyze the Zr-oxides dispersion phenomenon on  the glassy layer shown in Fig. 9 (i), the surface EDS mapping of the  20 wt.%ZrB2-80 wt.% SiC coating sample after oxidation are shown in  Fig. 10. Si, B and O elements were uniformly distributed in the glassy  layer, while Zr atoms were mostly concentrated in the precipitated  particles and dispersed in the glass only in negligible amount.  TEM analysis was undertaken to analyze the nature of the glassy  layer. As shown in Fig. 11 (a), besides some large ZrO2 particles, the  glass contained lots of smaller grains. A magnified image shown in  Fig. 11 (b) revealed that the small grains with regular shape were  nanocrystal particles about 20 nm in diameter, whose high-resolution  images  is  shown  in Fig. 11  (c). The nanocrystal particles were  embedded into the amorphous SiO2. The spherical nanocrystals were   Fig. 13. Sketch of  coating samples.   the oxidation protective mechanism of   the ZrB2-SiC   weakening. In addition, the intensity of the diffraction peak of ZrSiO4  phase gradually increased with the extension of oxidation time, indicating the generation of thermal stable ZrSiO4 phase by the Eq. (7).  However, due to the reduced ZrO2, the diffraction peak intensity of  ZrSiO4 phase did not continuously increase with further prolonging of  the oxidation time.  SEM images of the 80 wt.%ZrB2-20 wt.%SiC coating oxidized at  1973 K under different oxidation times in air are shown in Fig. 9. After  oxidation, the coating was obviously covered with glassy layer with  numerous white particles and pores. The pores were caused by the  release of CO, CO2, SiO or B2O3. After EDS point analysis, the white  particles could be determined as Zr-oxides. As shown in Fig. 9 (a) and (f),  after oxidation at 1973 K for 20 min, only a few ZrO2 particles were   JournaloftheEuropeanCeramicSociety41(2021)1059-10701067\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        primarily composed of Zr-oxides. The oxide dispersion mechanism is  illustrated in Fig. 11 (d). Since the oxidation temperature was higher  than the softening temperature of the amorphous SiO2 glass, the SiO2  glass would form during oxidation, in which Zr-oxides had very little  solubility and would soon precipitate. Moreover, under the action of  surface  tension, the molten SiO2 glass with high viscosity would   Fig. 15. (a) The schematic diagram of the glassy layer to prevent oxygen  penetration; (b) oxygen blocking mechanism of the Zr-B-Si-O glassy layer.   gradually spread on the surface of the coating, making the embedded Zr oxides stripped away with the action of viscous force and dispersed in  the glass to form the Zr-B-Si-O glassy layer.  Fig. 12 (a)-(d) and (e)-(h) show the macro-morphology and backscatter SEM micrographs of the samples after oxidation, respectively. As  show in Fig. 12 (a), when the content of ZrB2 was 20 wt.%, the surface of  glassy layer was rough with majority of large pores and less zirconia.  With increasing the ZrB2 content, the amount of Zr-oxides increased as  well, while the surfaces got smoother with lesser and smaller pores.  Therefore, with increasing the dispersed Zr-oxides nanocrystal particles  in SiO2 glass, the stability and protective effects of the Zr-B-Si-O glassy  layer at 1973 K in air were strengthened. Fig. 12 (i)-(l) show the cross section EDS oxygen mapping in ZrB2-SiC coatings with different ZrB2  content oxidized at 1973 K in air for 100 min. With increasing the ZrB2  content, the thickness of the glassy layer on the coated samples resulted  740 μm, 660 μm, 296 μm and 185 μm, respectively. The obviously  decreased  thickness of  the glassy  layer  further demonstrated  the  enhancement of the deactivation ability of the coatings as raising the  ZrB2 content.  Fig. 13 shows the sketch of the oxidation protective mechanism of  the ZrB2-SiC coating samples during oxidation. The oxidation resistance  of the ZrB2-SiC coating derived first from the capability of the glassy  layer to slow down oxygen penetration and then from the physical  barrier of the coating itself, under whose cooperation the permeation of  oxygen into the carbon substrate was greatly inhibited.   3.4. Oxygen blocking mechanisms of the coatings at 1973 K   Since the mass loss of the carbon substrate in the active oxidation  stage was an effective parameter to evaluate the structural blocking  ability of the coating itself, we established the concept of the structure  factor (αs,k), shown in Eq. (12).   Fig. 14. Structure factor (a) and deactivation factor (b) curves of the coatings;  (c) oxygen permeability of the ZrB2-SiC coating samples in the steady oxidation  stage. For definition of these parameters, see Eqs. (12)-(14), respectively.   αs,k =  1  n  [(cid:0)  ∑n  ms,k,j  (cid:0)  j=1  1  ms,k,j(cid:0) mp,j(cid:0)  (cid:0)  )  (cid:0)  (cid:0)  1 (cid:0)  (cid:0)  mc,k,j mp,j  )  ) ]  mc,k,j(cid:0)  1  (12)    Where αs,k is the structure factor of k formulation sample; mc,k,j and ms,k,j  are weight per unit area of k formulation coating and bulk reference   JournaloftheEuropeanCeramicSociety41(2021)1059-10701068\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        samples at j time respectively; mp,j is the weight per unit area of pure  graphite at j time. The subscript s, c and p indicate bulk sample, coating  sample and graphite substrate, respectively.  The smaller the αs,k is, the better the coatings protect the substrate.  As shown in Fig. 14 (a), with raising the ZrB2 content, the αs,k value of  the coatings gradually decreased. Thus, the raised ZrB2 content in  coatings enhanced the structural blocking ability of the coatings themselves, which were caused by the gradually denser structures shown in  Fig. 2 (f). When the glassy layer with self-sealing characteristic was  completely formed, the effect of the glass was mainly related to the  chemistry of the glass. To assess the blocking ability of the glassy layer  for oxygen diffusion, the bulk reference samples were chosen as the  research objects to eliminate the interference of substrate oxidation on  the oxidation curves of the coating samples, based on which the concept  of the deactivation factor (αi,k) was established, as shown in Eq. (13).  The smaller the αi,k is, the better the inhibition effect of the glassy layer  on the diffusion of oxygen is. As shown in Fig. 14 (b), with increasing  ZrB2 content,  the αi,k value of  the samples gradually decreased,  demonstrating that the raised ZrB2 content in the coatings enhanced the  blocking ability of the glassy layer.   αi,k =  1  n  (cid:0)  ∑n  j=1  ms,k,j+2 (cid:0)  ms,k,j+1  )  tj  (13)    Where αi,k is the deactivation factor of k formulation sample; ms,k,j+1 is  weight per unit area of k formulation bulk reference samples at j time  respectively; tj is time interval of weighing (tj = 10 min).  To disclose the barrier effect of the glassy layer on the diffusion and  permeation of oxygen to the carbon substrate, based on Eq. (14), we  detected the oxygen permeability of the ZrB2-SiC coating samples in the  steady oxidation stage, shown in Fig. 14 (c).   (cid:0)  ms,k,j  (cid:0)  Pj =  )  ms,k,0 mp,j  (cid:0)  (cid:0)  (cid:0)  mc,k,j  (cid:0)  mc,k,0  )  mp,0  × 100%  (14)    Where Pj is the oxygen permeation at j time; ms,k,j , and mc,k,j are weight  per unit area of k formulation bulk reference and coating samples in j  time interval; mp,j is the weight per unit area of pure graphite in j time  interval. ms,k,0 , mc,k,0 and mp,0 are initial weight of the samples. The  subscript s, c and p indicate bulk reference sample, coating sample and  graphite substrate, respectively.  With extending oxidation time, the oxygen permeability curves of  the ZrB2-SiC coatings containing less ZrB2 content present a steady trend  (20 wt.%) or a gradual upward trend (40 wt.%), while that containing  more ZrB2 content (>40 wt.%) exhibited a downward trend, indicating  the high ZrB2 content conduced to inhibit the oxygen permeability at  1973 K. As shown in Fig. 12 (e) (f), due to the release of gas by products, the generated bubble holes in the glassy layer provided  channels for the diffusion of oxygen, which led to the decline of the  protective ability of the glassy layer, thus resulting in the upward trend  of the oxygen permeability curve in the steady oxidation stage. In  addition, with increasing the ZrB2 content, as shown in Fig. 12 (g)-(h),  the stability and integrity of the Zr-B-Si-O glassy layer were significantly  enhanced, which were responsible for the evidently decreased oxygen  permeability and αi,k value of the ZrB2-SiC coatings in the steady  oxidation stage.  As the most vital barrier layer between the coatings and oxygen, the  glassy layer acted as a barrier to the penetration of oxygen. The schematic diagram that shows how the glassy layer prevents oxygen penetration is shown in Fig. 15 (a). However, after the formation of the glassy  layer, the release of gas by-products, such as CO, CO2, SiO or B2O3,  would form oxidation pores on the glass film and the coatings [35].  Although the outside glassy layer had self-sealing effect during oxidation, the inside oxide holes could not be filled and sealed promptly.  Moreover, the release of by-products and the existence of oxide holes  would form negative pressure inside the coatings, while oxygen outside   the coatings applies positive pressure. Hence, under the influence of  pressure difference, the penetration of oxygen into the glassy layer  would be aggravated, thus weakening the stability of the glassy layer.  Therefore, with increasing the number of internal defects in the coatings, the structural factor αs,k of the coatings increased, the permeability  of oxygen increased, and the structural blocking ability of the coatings  became worse.  Fig. 15 (b) shows the oxygen blocking mechanism of the Zr-B-Si-O  glassy layer. As mentioned above, the dispersion of Zr-oxides nanocrystal particles in the SiO2 glassy layer was the essential reason of the  formation of the Zr-B-Si-O glassy layer. Hence, due to the higher melting  point of the Zr-oxides, the un-melted Zr-oxides nanocrystal particles  actually existed in the fluid molten glass as heterogeneous hard particles, thereby enhancing the overall un-melted area of the glassy layer.  Since the diffusion resistance of the oxygen in the un-melted region was  much larger than that of the fluid molten glass, the un-melted Zr-oxides  would force the oxygen to turn round or deflect in the diffusion process.  Besides that, since the effusion of gas by-products and the permeation of  oxygen both passed through the oxygen diffusion channels in the glassy  layer, which would form pressure stress on the oxygen diffusion channels. Owing to the high valence state and small radius of the Zr4+ dispersed in the glassy layer, the Zr4+ was apt to interact with [SiO4]  tetrahedral units, forming huge complex network structures, increasing  the overall viscosity of the glassy layers. With the increase of the viscosity of the molten glass, the enhanced viscous force of the glassy layer  would produce a repulsive force to resist the pressure stress caused by  oxygen permeation and effusion of gas by-products. Therefore, the oxygen diffusion channels would be narrowed and decreased,  thus  significantly restraining the diffusion speed of oxygen through the glassy  layer, and weakening the damage of oxygen to the glassy layer. Therefore, with increasing the content of Zr-oxides, the increased viscosity  enhanced the viscous force of the glassy layer, making the oxygen  permeability of the glassy layer gradually decreased, as shown in Fig. 14.  However, if the content of Zr-oxides exceeded a certain degree, the  excessive viscosity of glassy layer would reduce its self-healing ability  and increase the oxygen permeability of the glassy layer instead, thereby  weakening the oxidation protection efficiency of the coatings.   4. Conclusion   In this paper, ZrB2-SiC coatings with different ZrB2 content ranging  from 20 to 80 wt.% were prepared on graphite substrate by SPS technology. The oxidation behavior of the various coating was studied in a  static furnace at 1700  C up to 100 min and compared with that of the  corresponding bulk ZrB2-SiC references. With increasing ZrB2 content,  the maximum oxygen permeability of the coating at 1973 K decreased  from 2.91 % to 1.77 %, while the oxidation protection efficiency  increased from 96.4%-97.7%.  In the first oxidation stage up to 30 min, although increasing ZrB2  content activated the oxidation reaction of the coatings, the damage to  the coating caused by the SiC active oxidation was inhibited, thus  enhancing the structural blocking ability of the coatings.  In the second oxidation stage from 40 min to 100 min, the dispersion  of Zr-oxides nanocrystal particles in SiO2 glass enhanced the stability  and protective ability of the glass, which intensified with extending  oxidation time.  High precipitation of Zr-oxides nanocrystal particles conduced to  broaden crystalline areas and increased the viscosity of the glassy layer,  so that the glassy layer on the 60ZrB2-40SiC coating exhibited the best  oxygen blocking ability.  Due to the enhanced structural blocking ability of the coatings  themselves and the blocking ability of the glassy layer formed on top of  them, increasing ZrB2 content enhanced the oxidation stability of the  coating at 1973 K up to 100 min, making the oxygen permeability of the  coating gradually decrease.   JournaloftheEuropeanCeramicSociety41(2021)1059-10701069\\x0c', 'X. Ren et al.                                                                                                                                                                                                                                        Declaration of Competing Interest   The authors declare that they have no known competing financial  interests or personal relationships that could have appeared to influence  the work reported in this paper.   Acknowledgement   This work has been supported by the National Natural Science  Foundation of China (Grant No. 51972338, 51874305), National Defense Basic Research Program (JCKYS2019607004-01), Open Sharing  Fund for the Large-scale Instruments and Equipments of CUMT (DYGX 008, DYGX-009, DYGX-010). The Authors acknowledge the support of  TEM tests and analyses provided by the Advanced Analysis & Computation Center (China University of Mining and Technology).   Appendix A.  Supplementary data   Supplementary material related to this article can be found, in the  online version, at doi:https://doi.org/10.1016/j.jeurceramsoc.2020.10  .036.   References   [5]  [1]  J.C. Ren, Y.L. Zhang, P.F. Zhang, T. Li, J.H. Li, Y. Yang, Ablation resistance of HfC  coating reinforced by HfC nanowires in cyclic ablation environment, J. Eur. Ceram.  Soc. 37 (2017) 2759-2768.  [2] H.J. Li, H.J. Luo, L. Li, L.H. 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},{
  "_id": 57,
  "PDF": "Effect of thermal exposure on strength of ZrB2-based composites with nano-sized SiC particles.pdf",
  "Text": "['Composites Science and Technology 68 (2008) 3033-3040  Contents lists available at ScienceDirect  Composites Science and Technology  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o m p s c i t e c h  Effect of thermal exposure on strength of ZrB2-based composites with nano-sized SiC particles  Shu-Qi Guo a,*,  Jenn-Ming Yang b, Hidehiko Tanaka c, Yutaka Kagawa a,d  a Composites and Coatings Center, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, b Department of Materials Science and Engineering, University of California, Los Angeles, CA 90095-1595, USA c Nano Ceramic Center, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan d Research Center for Advanced Science and Technology, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8505,  Japan  Japan  a r t i c l e  i n f o  a b s t r a c t  The effect of oxidation exposure on the room-temperature flexural strength of ZrB2-based composites with nano-sized and micron-sized SiC particles was investigated. The composites were densified by vacuum hot-pressing at 2000 °C for 60 min under a uniaxial pressure of 30 MPa. The flexural strength of asreceived composites and after exposure in air at 1400 °C for 10 h was measured using the four-point bending test. The results show that the presence of intergranular and intragranular nano-sized SiC particles imparts a better oxidation resistance and improves the flexural strength of single-phase ZrB2 and ZrB2 ceramics with micron-sized SiC.  Ó 2008 Elsevier Ltd. All rights reserved.  Article history:  Received 26 March 2008 Received in revised form 12 June 2008 Accepted 17 June 2008 Available online 1 July 2008  Keywords:  Zirconium diboride Nano SiC particles Thermal exposure Flexural strength  1. Introduction  Diborides and carbides of zirconium (ZrB2 and ZrC) have extremely high melting points (>3000 °C), high thermal and electrical conductivities, chemical inertness against molten metals, and great thermal shock resistance [1,2]. The unique combinations of mechanical and physical properties make them attractive candidates for structural applications at ultra-high temperatures. As a result, ZrB2 and ZrC ceramics are being considered for a variety of high-temperature thermomechanical structural applications, including furnace elements, plasma-arc electrodes, rocket engines and thermal protection structures for leading-edge parts on hypersonic re-entry space vehicles at over 1800 °C [1-6]. However, the use of those single-phase materials for high temperature structural applications is limited by the poor oxidation and ablation resistance as well as poor damage tolerance. The composite approach has been successfully adopted in order to improve the oxidation and ablation resistance of single-phase ceramics. For example, the addition of a second phase, such as SiC to ZrB2 results in a composite with improved strength, better oxidation, thermal shock and ablation resistance [2,7-12]. The improvement of oxidation and ablation resistance is believed to arise from the formation of a coherent passivating oxide scale on the surface. Furthermore, the addition of SiC also limits the grain  * Corresponding author. Tel.: +81 (0)29 859 2223; fax: +81 (0)29 859 2401. E-mail address: GUO.Shuqi@nims.go.jp (S.-Q. Guo).  0266-3538/$ see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2008.06.021  growth during densification and improves the sinterability of ZrB2. Thus far, micron-sized and ultra-fine SiC particles were added to the ZrB2 for most of the studies reported in the literature [9,10]. The SiC particles tend to be dispersed along the grain boundaries of ZrB2 (intergranular phase). It had been recently demonstrated that some SiC grain growth occurred during hot-pressing of ZrB2-SiC composites [13,14]. Furthermore, the flexural strength of the composite decreased substantially as the average size of SiC grains increased from \\x181.2 to 3.1 lm. Zhu et al. [14] also suggested that the largest SiC grains in the microstructure acted as the critical flaws causing the failure of the composite. Hence, it was concluded that smaller SiC grain sizes would result in higher strength for ZrB2-SiC composites. Very recently, Hwang et al. [15] reported that the incorporation of nano-sized SiC particles improved the oxidation resistance of ZrB2 ceramics. However, the retained fracture strength of nano-sized SiC-ZrB2 ceramics after oxidation exposure at high temperature is not well-known. In the present study, the effects of oxidation exposure on the room-temperature flexural strength of hot-pressed ZrB2 composites with nano-sized and micron-sized SiC particles were examined in air at 1400 °C, for up to 10 h.  2. Experimental procedure  The starting powders used in this study were: ZrB2 powder (Grade F, Japan New Metals, Tokyo), average particle size \\x192.1 lm, nano b-SiC powder (Sumitomo Osaka Cement Co. Ltd.,  \\x0c', '3034  S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  Japan), average particle size \\x1930 nm and oxygen of 0.38 Osaka, wt.%; and a-SiC powder (GC#2000, Showa Denko K.K, Tokyo, Japan), average particle size \\x196.4 lm and oxygen of 0.29 wt.%. In order to examine the effects of nano-particle content as well as starting particle size on strength, five batches of powder were prepared, four of them containing nano b-SiC powder of 5, 10, 15 and 20 vol.%, respectively, the fifth one containing 20 vol.% large a-SiC powder. In addition, the single-phase ZrB2 powder was prepared to compare the effect of SiC addition on flexural strength. Compositions of the ZrB2-based composites with SiC are listed in Table 1. The powder mixtures were ball-milled using SiC milling media and ethanol under 200 rpm for 24 h, and the resulting slurry was then dried under magnetic stirring to avoid sedimentation. Before sintering, the dried mixtures were sieved through a metallic sieve with \\x0060-mesh screen size. The obtained powder mixtures were hot-pressed (FVHP-1-3, Fuji Electric Co. Ltd., Tokyo, Japan) in the graphite dies under a pressure of 30 MPa under vacuum in tablets 21 mm \\x02 25 mm \\x02 3.5 mm in size. Powder compacts were heated under vacuum of \\x187.0 \\x02 10\\x003 Pa to 2000 °C with a heating rate of 50 °C/min under an uniaxial pressure of 30 MPa. After hot-pressing at 2000 °C for 60 min, the temperature was decreased to 500 °C with a heating rate of 20 °C/min, and then the sample was cooled to 25 °C (room temperature) in the furnace. The load was removed when the die temperature dropped below 1750 °C. The densities, q, of the hot-pressed composite compacts were measured by the Archimedes method with distilled water as the medium. The grain size, d, was determined by measuring the average linear intercept length, dm, of the grains in FE-SEM images of sintered composites, according to the relationship of d = 1.56dm which was given by Mendelson [16]. Test specimens with dimensions of 25 mm \\x02 2.5 mm \\x02 2 mm were cut from the sintered ZrB2-SiC composites plates. All the specimens were ground with an 800-grit diamond wheel, and the tensile surfaces for the bend test were polished with diamond paste up to 0.5 lm. The edges of the specimen were then chamfered at 45°. In addition, in order to examine strength retention of the sintered materials after exposure in air at a high temperature, the specimens were exposed in air at 1400 °C. The surfaces of the specimens were ultrasonically cleaned in acetone and kept in an oven with a constant temperature of 100 °C prior to oxidation exposure. The specimens were exposed to dry air at 1400 °C, for up to 10 h. The 1400 °C oxidation exposures were conducted in an electronic furnace (Model SSFT-1520, Nikkato Co. Ltd., Tokyo, Japan). The specimens were supported on high purity Al2O3 knifeedged fixtures during oxidation exposure. The fracture strengths of the as-sintered and post-oxidized specimens were determined using the four-point flexure test at room temperature (inner span 10 mm, outer span 20 mm). The bending test was performed using an universal testing system (Autograph Model AG-50KNI, Shimadzu Co. Ltd., Kyoto, Japan) with a constant crosshead speed of 0.5 mm/min. At least three specimens were used for each measurement. After the bending test, the fracture surfaces of the specimens were examined by  field emission scanning electron microscopy Hitachi High-Technologies Corporation, Tokyo,  (FE-SEM, Japan).  S-4800,  3. Results and discussion  3.1. Microstructures  Typical microstructural features of the single-phase ZrB2 and ZrB2-based composites with nanoand micron-sized SiC observed under FE-SEM are shown in Fig. 1. The microstructure of the singlephase ZrB2 consists of the equiaxed ZrB2 and some pores (dark contrast). Most of the pores are presented at multiple-grain pockts. On the other hand, the general microstructures of the ZrB2-based composites with nanoand micron-sized SiC are similar, consisting of the equiaxed ZrB2 (grey contrast) and SiC (dark contrast) grains. The measured average grain sizes of the ZrB2 and SiC are listed in Table 1. In the case of the single-phase ZrB2, the average grain size of ZrB2 is \\x186.1 lm, which is significantly higher than the starting (\\x182.1 lm). powder Apparently, grain growth occurred during hot-pressing. The average grain size of ZrB2 in the composite with 5 vol.% nano-sized SiC particle is similar to that in the single-phase ZrB2. This means that the 5 vol.% nano-sized SiC addition is insufficient for hindering the grain growth of ZrB2 phase during sintering. However, the average grain size of ZrB2 in the composites with 10-20 vol.% nano-sized SiC particle is in the range of 4.5-4.2 lm which is smaller that in the single-phase ZrB2. This indicated that 10 vol.% or more SiC addition is sufficient for hindering the grain growth of ZrB2 during the sintering. In addition, for the ZrB2-based composite with micron-sized SiC, the grain size of ZrB2 is measured to be 5.4 lm which is larger that of the composite with nano-sized SiC particle. This implies that the nano-sized SiC is more effective for hindering the grain growth of ZrB2 than micro-sized SiC. The average grain size of the coarse SiC was reduced from 6.0 to 1.6 lm after sintering. Apparently, the mixing process is effective in reducing the size of SiC particles. However, the average size of SiC in the ZrB2 composites with nano-sized SiC is \\x180.8 lm, which is substantially higher than that of the starting powder (\\x1830 nm). The mixing process employed was not effective in breaking apart the agglomeration of nano-sized SiC particles. As a result, clusters of nano-sized SiC particles were fused together during hot-pressing to form SiC particles with submicron in diameter. Occasionally, several nano-sized SiC particles can be observed within the ZrB2 grains. As a result, a more effective mixing/dispersion technique to break apart the agglomeration nano-sized SiC particles needs to be developed to uniformly disperse the nano-sized SiC particles. The measured densities and relative densities for the various ZrB2-SiC composite materials are also summarized in Table 1. It that single-phase ZrB2 has the lowest density (\\x1891%). is evident The low relative density is expected due to its strong covalent bond and low self diffusion [11]. Melendez-Martinez et al. [17] reported a final density of 86.5% for a single-phase ZrB2 densified under a pressure of 30 MPa at 1900 °C for 30 min. Higher density can be achieved through optimizing the temperature, pressure, time, additives and particle sizes of the staring materials. In the present  Table 1 Compositions, true densities, relative densities and grain size of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles  Materials  Compositions (vol.%)  SiC grade  Theoretical density (g/cm3)  Measured density (g/cm3)  Relative density (%TD)  Average grain size (lm)  ZSN-0 ZSN-5 ZSN-10 ZSN-15 ZSN-20 ZSM-20  ZrB2  100 95 90 85 80 80  SiC  0 5 10 15 20 20  - Nano Nano Nano Nano Micron  6.09 5.95 5.80 5.66 5.52 5.52  5.51 5.69 5.65 5.46 5.37 5.42  90.4 95.6 97.4 96.5 97.3 98.2  ZrB2  6.1 ± 2.2 5.9 ± 1.4 4.5 ± 1.6 4.4 ± 1.7 4.2 ± 1.9 5.4 ± 1.6  SiC  - 0.8 ± 0.4 0.8 ± 0.4 0.9 ± 0.5 0.9 ± 0.5 1.6 ± 0.6  \\x0c', 'S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  3035  Fig. 1. Typical FE-SEM micrographs of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles; (a) ZSN-0, (b) ZSN-10, (c) ZSN-20 and (d) ZSM-20.  study, the relative densities exceeding 95% are obtained for the ZrB2 powder with nano-sized particles, which is significantly higher than that of the single-phase ZrB2 powder. It is evident that the addition of nano-sized SiC has a similar effect in improving the sinterability of ZrB2. Improvement of the sinterability of ZrB2 due to ultra-fine SiC addition is documented in literatures. Monteverde [9] showed that ZrB2 with 10 vol.% ultra-fine SiC (d90 = 0.8 lm) achieved full density by vacuum hot-pressing at 1900 °C for 20 min. Hwang et al. [15] showed that the sinterability of ZrB2 was significantly improved by addition of nano-sized SiC and the improvement was enhanced with SiC amount as well as with reducing SiC starting-powder size. Zhu et al. [14] also showed that smaller starting SiC particle sizes led to improved densification. This improvement of densification is due to the formation of intergranular liquid phases during hot-pressing, assisting in densification at lower temperatures. Similar cause is expected in the present studied ZrB2 with nanoand micron-sized SiC starting powders. However, the highest relative density is not obtained in the ZrB2 with nano-sized SiC particle, but in the ZrB2 with micron-sized SiC particle. This probably is associated with the agglomeration nano-sized SiC particles which lead to the formation of submicron SiC (Table 1), in turn resulting in reduced improvement in densification of ZrB2 with nano-sized SiC particles.  3.2. Oxidation  Fig. 2 shows typical examples of X-ray diffraction patterns of the single-phase ZrB2 and ZrB2-based composites with nanoand micron-sized SiC before and after 10 h of oxidation exposure at 1400 °C. Before oxidation exposure (Fig. 2a), only ZrB2 phase was detected in the single-phase ZrB2, while ZrB2 and SiC phases were present in the ZrB2 with nanoand micron-sized SiC. Although a grain-boundary phase is reported in ZrB2-based composites with SiC particles [15], a trace amount of the secondary crystalline phase was not detected in any instance. In addition, the peak intensities of ZrB2 phase decreased with SiC addition, and the intensities were stronger for the ZrB2 with micron-sized SiC than for that with nano-sized SiC. This indicated that the amount of SiC particles in per unit area is higher in the ZrB2 with nano-sized SiC than in that  with micron-sized SiC. After oxidation exposure (Fig. 2b), only ZrO2 phase was detected for all the compositions materials. The peak intensities of the ZrO2 phase were the strongest in the single-phase ZrB2. Furthermore, the intensities were stronger for the ZrB2 with micron-sized SiC than for the ZrB2 with nano-sized SiC as well as they weakened with amount of nano-sized SiC particles. This implies that addition of SiC particles improved oxidation resistance of the single-phase ZrB2. This improvement was better in the ZrB2 ceramics with nano-sized SiC than in the ZrB2 with micronsized SiC and it was enhanced with amount of nano-sized SiC. These XRD analyses indicated that the single-phase ZrB2 and ZrB2-containing nanoand micron-sized SiC particles were oxidized to form the ZrO2 phase on the surfaces of specimens when they were exposed to 1400 °C in dry air. Fig. 3 showed the appearances of the hot-pressed ZrB2-SiC composites before and after the oxidation exposure. Before oxidation exposure, all the specimens were almost the same color. After oxidation exposure, however, the color of single-phase ZrB2 turned white with clear cracking on the surface. At high-magnification image (Fig. 4a), no distinct difference was observed on the oxidized surface. This shows that the oxidized surface consists of singlephase ZrO2. It is well-known that the single-phase ZrB2 was oxidized to form ZrO2 and B2O3 when ZrB2 was exposed to elevated temperature in air [7,18]. Thus, in the present study, when the single-phase ZrB2 was exposed to air at 1400 °C, the ZrB2 was oxidized to the ZrO2 and B2O3 phases according to following reaction  ZrB2 þ 5 2  O2 ! ZrO2 þ B2O3 :  ð1Þ  Although the B2O3 liquid layer presents an effective barrier to the transport of oxygen below 1100 °C, at 1400 °C and above, the rate of vaporization of B2O3 is comparable to the rate of formation [7]. Thus, only ZrO2 was presented for the single-phase ZrB2 after oxidation exposure. In addition, some cracks were observed on the oxidized surface (indicated by arrows in Fig. 4a). Cracking was induced by the volume expansion of the ZrO2 phase (typically 3-5 vol.%), which was produced during exposure through the phase transformation of tetragonal to the monoclinic phase upon cooling from exposure temperature to room temperature.  \\x0c', '3036  S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  )  .  u  .  a  (  y  t  i  s  n e  t  n  I  ZrB2  SiC  ZSN-0  ZSM-20  ZSN-20  )  .  u  .  a  (  y  t  i  s  n e  t  n  I  ZrO2  ZSN-0  ZSM-20  ZSN-20  20  30  40  50  60  70  80  2θ(degree)  20  30  40  50  60  70  80  2θ(degree)  Fig. 2. Typical examples of X-ray diffraction patterns of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles; (a) before and (b) after oxidation exposure.  Fig. 3. Typical macroscopic appearances of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles before and after oxidation exposure.  On the other hand, the color of ZrB2-based composites with nanoand micron-sized SiC particles turned black (Fig. 3b), with a transparent glassy film on the surfaces of the oxidized specimens. In particular, the color of the ZrB2-based composites with nanosized SiC particles is darker than that of the ZrB2-based composites with micron-sized SiC, which is a clear indication that more SiC and ZrB2 particles remained intact (without being oxidized) after oxidation exposure. Under SEM image (Fig. 4b), the oxidized surface of ZrB2 with SiC showed a distinctly difference with the sin gle-phase ZrB2: some white particles embedded in the background. The EDX analysis of the oxidized surfaces showed that the light places were rich in Zr and the background was rich in Si, with a high concentration of O (Fig. 5). The EDX results indicated that the oxidized surface of ZrB2-based composites with SiC consists of silica glass background and some ZrO2 particles embedded it. It is known that the ZrB2 and SiC were oxidized simultaneously to form ZrO2, B2O3 and silica glass when the ZrB2-containing SiC composition was exposed to air at high temperature [7,19]. Our      \\x0c', 'S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  3037  Fig. 4. SEM images of oxidized surfaces of the hot-pressed single-phase ZrB2 and ZrB2-based composite with 20 vol.% nano-sized SiC particles; (a) ZSN-0 and (b) ZSN20.  experimental results indicated that the oxidation of ZrB2 and SiC might be occurring for the studied material during the oxidation exposure. Moreover, the weight gain after oxidation in the ZrB2 with nano-sized SiC particles decreased with increasing amount of SiC particle (Table 2). Furthermore, the weight gain after oxidation exposure in the ZrB2-based composite with 20 vol% nano-sized SiC particles is lower than that with 20 vol% micron-sized SiC particles. It is evident that the ZrB2-based composites with nano-sized SiC particles exhibited a better oxidation resistance than that with micron-sized SiC particles. Better oxidation resistance due to addition of nano-sized SiC particles is documented in literature. Hwang et al. [15] revealed that SiC grain-size reduction result in improved oxidation resistance in ZrB2-based ceramics with SiC. This improvement is attributed the formation of the protective silicarich glass layer on the surface of the ZrB2-based composites with nano-sized SiC during the earlier stage of oxidation, as a result of an increase in the ZrB2/SiC interface length per unit area of exposed-surface and a decrease in the spacing between SiC particles with increasing amount of SiC particles or reducing SiC starting size. Similar cause is expected in the nano-sized SiC-containing ZrB2 ceramics investigated in this study.  3.3. Flexural strength  Table 2 also lists the flexural strengths of the single-phase ZrB2 and ZrB2-based composites with nanoand micron-sized SiC particles measured by four-point bending test before and after oxida Fig. 5. Elemental composition maps of oxidized surface of based composite with 20 vol.% nano-sized SiC particles.  the hot-pressed ZrB2 Table 2 Weight gain and flexural strength of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles before and after 10 h of oxidation exposure at 1400 °C in dry air  Materials  Total weight gain (mg/cm2)  4-Point flexural  Strength (MPa)  ZSN-0 ZSN-5 ZSN-10 ZSN-15 ZSN-20 ZSM-20  24.01 7.11 6.94 5.31 4.66 11.02  As-received  Post-exposed  457 ± 58 549 ± 49 524 ± 63 714 ± 59 608 ± 93 531 ± 10  141 ± 21 635 ± 15 610 ± 83 718 ± 82 700 ± 41 506 ± 19  \\x0c', '3038  S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  tion of 10 h at 1400 °C. Before oxidation exposure, the single-phase ZrB2 has the lowest flexural strength (457 MPa). This may be due to the presence of \\x1810% porosity. The flexural strength of the ZrB2based composites with nano-sized SiC particles is significantly higher than that of the single-phase ZrB2. The ZrB2-based composite with 15 vol.% nano-sized SiC has the highest flexural strength of 714 MPa. It is also clear that the flexural strength of ZrB2-based composite with 20 vol.% nano-sized SiC is higher than that with 20 vol.% micron-sized SiC. This verifies an earlier study that smaller SiC grain sizes would result in higher strength for ZrB2-SiC composites [14]. After oxidation exposure, the flexural strength of single-phase ZrB2 dropped sharply from 457 MPa to 141 MPa, which corresponds to \\x1870% reduction. In contrast, the flexural strength of ZrB2-based composites with nano-sized SiC particles increased significantly except for the composite with 15 vol.% SiC in which its flexural strength retained almost constant. The flexural strength of the ZrB2-based composite with micron-sized SiC particles decreased after oxidation, however. The fracture surfaces of the hot-pressed ZrB2 and ZrB2-based composites with nanoand micron-sized SiC are observed under SEM, which examples are shown in Fig. 6. A typical intergranular fracture mode was observed in the single-phase ZrB2 as shown in Fig. 6a. For the ZrB2-based composites with nano-sized SiC particles (Fig. 6b), both ZrB2 and SiC grains exhibited typical intragranular fracture characteristics. SiC particles were located along the grain boundaries (intergranular). In addition, high-magnification SEM images shown in Fig. 6c revealed that some nano SiC particles were observed within ZrB2 grains (intragranular) as indicated by arrows in Fig. 6c. Niihara [20] reported that the strength of a ceramic could be enhanced substantially through the incorporation of nano-particles within the grains (intragranular nanocomposites). For the ZrB2-based composite with micron-sized SiC (Fig. 6d), the fracture surface of ZrB2 and SiC also exhibited a typical intragranular feature. The large SiC particles were observed at the multipleZrB2 grains pockets as indicated by arrows in Fig. 6d. As a result, the nano-sized SiC particles are more effective in enhancing the strength of the ZrB2 ceramic than that derived from micron-sized SiC particles.  The macroscopic fracture appearances and SEM micrographs of the fracture surfaces for the hot-pressed ZrB2-based composites with nanoand micron-sized SiC after being oxidized at 1400 °C for 10 h are shown in Fig. 7. For single-phase ZrB2, only central part of specimen was not oxidized after the oxidation exposure (Fig. 7a), and two large cracks were extended through to near the unreacted ZrB2 (indicated by arrows). The large cracks led to a substantial reduction of the strength of ZrB2 after oxidation exposure. Under high-magnification SEM image (Fig. 7d), the unreacted ZrB2 showed typical intragranular fracture characteristics. This differed from the as-sintered ZrB2 which showed a typical intergranular fracture mode (Fig. 6a). For both the ZrB2-based composites with nanoand micronsized SiC (Fig. 7b and c), on the other hand, the observable oxidized reacted region are absent, as a result of improved oxidation resistance of ZrB2 through the addition of SiC particles. Although, for the ZrB2 with micro-sized SiC, the fracture originated from a defect at or near the surface (indicated by an arrow in Fig. 7b), the fracture origin was not distinguishable for the ZrB2 with nano-sized SiC (Fig. 7c). Higher-magnification SEM images, taken from the tensile surface of specimens, are shown in Fig. 7e and f, respectively. For the oxidized ZrB2-based composite with micron-sized SiC particles (Fig. 7e), a thicker glassy layer is observed and a defect was also observed at the interface between the oxide scale and bulk ZrB2 (indicated by an arrow). This revealed that the ZrB2 specimens with micron-sized SiC develop thick oxide scales that probably possess high level of defects, resulting in a lower strength after oxidation exposure. In contrast, for the oxidized ZrB2-based composites with nano-sized SiC particles (Fig. 7f), the thinner oxide scale was observed. The thin oxide scale of the ZrB2-based composites with nano-sized SiC is attributed to its better oxidation resistance than that with micron-sized SiC. Clearly, the formation of the thin oxide scale can heal the surface flaws without creating new cracks and defects at the oxidized surface and play a significant role in the strength improvement after oxidation exposure. Thus, the improvement of flexural strength after oxidation exposure may be attributed to the formation of a thinner oxide scale which effectively alters the geometry of the surface defects [21], as a result of a  Fig. 6. Typical fracture surfaces of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles; (a) ZSN-0, (b) ZSN-15, (c) high-magnification of (b) and (d) ZSM-20.  \\x0c', 'S.-Q. Guo et al. / Composites Science and Technology 68 (2008) 3033-3040  3039  Fig. 7. Typical macroand microscopic fracture appearances of the hot-pressed ZrB2-based composites with nanoand micron-sized SiC particles after oxidation exposure; (a), (d) ZSN-0, (b), (e) ZSM-20, (c), (f) ZSN-20.  better oxidation resistance of that with micron-sized SiC.  the ZrB2 with nano-sized SiC than  4. Summaries  The room-temperature flexural strength of ZrB2-based composites with nano-sized SiC particles was investigated and compared to that with micron-sized SiC particles before and after oxidation exposure in dry air at 1400 °C for 10 h. The strength for ZrB2 with nano-sized SiC particles is higher than that with micron-sized SiC particles. This is attributed to a more uniform dispersion of nano-sized SiC particles at grain boundaries and within the grains, whereas the micron-sized SiC particles are located at multiplegrain pockets. After oxidation exposure, the flexural strength of ZrB2-based composites with nano-sized SiC particles increased. In particular, the flexural strength significantly increased for the composites containing 5, 10 and 20 vol.% nano-sized SiC particles. In contrast, the strength of the ZrB2-based composite with microsized SiC particles decreased after oxidation. The presence of intergranular and intragranular nano-sized SiC particles imparts a bet ter oxidation resistance and improves the flexural strength of single-phase ZrB2 and ZrB2-based ceramics with micron-sized SiC. Further improvement in strength may be achieved through a more uniform dispersion of nano-sized SiC particles.  Acknowledgements  J.M.Y. gratefully acknowledges the support from NASA Glen Research Center. Dr. N. Bansal is the program manager.  References  [1] Upadhya K, Yang JM, Hoffmann WP. Materials for ultrahigh temperature structural applications. Am Ceram Soc Bull 1997;76:51-6. [2] Fahrenholtz WG, Hilmas GE, Talmy IG, Zaykoski JA. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90:1347-64. [3] Brown AS. 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[10] Monteverde F, Bellosi A. Development and characterization of metal-diboridebased composites toughened with ultra-fine-SiC particles. Solid State Sci 2005;7:622-30. [11] Bellosi A, Monteverde F, Fabbriche DD, Melandri C. Microstructure mechanical properties of ZrB2-based Ceramics. J Mater Process Manuf 2000;9:156-70. [12] Chamberlain AL, Fahrenholtz WG, Hilmas GE. High-strength diboride-based ceramics. J Am Ceram Soc 2004;87:1170-2. [13] Rezaie A, Fahrenholtz WG, Hilmas GE. Effect of hot pressing time and temperature on the microstructure and mechanical properties of ZrB2-SiC. J Mater Sci 2007;42:2735-44.  zirconium  and Sci  and oxidationcontaining SiC  [14] Zhu S, Fahrenholtz WG, Hilmas GE. Influence of silicon carbide particle size on the microstructure and mechanical properties of zirconium diboride-silicon carbide ceramics. J Eur Ceram Soc 2007;27:2077-83. [15] Hwang SS, Vasiliev AL, Padture NP. Improved processing resistance of ZrB2 ultra-high temperature ceramics nanodispersoids. Mater Sci Eng 2007;A464:216-24. [16] Mendelson MI. Average grain size in polycrystalline ceramics. J Am Ceram Soc 1969;52:443-6. [17] Melendez-Martinez JJ, Dominguez-Rodriguez A, Monteverde F, Melandri C, de Portu G. Characterization and high temperature mechanical properties of zirconium boride-based materials. J Eur Ceram Soc 2002;22:2543-9. [18] Tripp WC, Graham HC. Thermogravimetric study of the oxidation of ZrB2 in the temperature range of 800 °C to 1500 °C. J Electrochem Soc 1971;118:1195-9. [19] Monteverde F, Bellosi A. Oxidation of ZrB2-based ceramics in dry air. J Electrochem Soc 2003;150:B552-9. [20] Niihara K. New design concept of structural ceramics-ceramic nanocomposites. J Ceram Soc Japan 1991;99:974-82. [21] Guo SQ, Hirosaki N, Yamamoto Y, Nishimura T, Mitomo M. Strength retention in silicon nitride ceramics with Lu2O3 additives after oxidation exposure in air at 1500 °C. 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},{
  "_id": 58,
  "PDF": "Effect of WB on oxidation behavior and microstructure evolution of ZrB2-SiC coating.pdf",
  "Text": "['Corrosion Science 155 (2019) 155-163  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Effect of WB on oxidation behavior and microstructure evolution of ZrB2-SiC coating  T  Chong Lia,b, Yaran Niua,⁎, Tao Liua,b, Liping Huanga, Xin Zhonga, Xuebin Zhenga,⁎, Chuanxian Dinga  a Key Laboratory of Inorganic Coating Materials CAS, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China b University of Chinese Academy of Sciences, Beijing 100049, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: ZrB2-SiC composite coating WB Vacuum plasma spray Oxidation behavior Microstructure  1.  Introduction  In this work, different contents (5, 10, 15 mol.%) of WB were introduced into ZrB2-SiC coatings which were fabricated by vacuum plasma spray technique. The oxidation resistant property of the ZrB2-SiC-WB ternary composite coatings was evaluated in static air at 1500 °C. The results showed that appropriate amount of WB addition significantly improved the oxidation resistance of ZrB2-SiC coating, as proved by the reduced weight gain and increased thickness of the liquid layer. The influence of WB and its oxidation product WO3 were analyzed based on the microstructure changes and thermodynamic calculations in detail.  With the key requirements for hypersonic vehicles to withstand high temperatures in oxidant atmosphere, it is necessary to develop outstanding thermal protection systems for extreme environment applications. Ultra-high temperature ceramics (UHTCs) have attracted great attention due to their excellent adaptability for critical conditions [1,2]. ZrB2 is one of the most important candidate materials in UHTCs family due to its high melting point (3050 °C), high thermal conductivity (˜100 W/(m K)) and relatively low density (6.09 g/cm3) [3,4]. The oxidation behavior of ZrB2 has been studied since 1960s [5,6]. It has been found that ZrB2 is easy to be oxidized above 800 °C, forming B2O3 and ZrO2. The formed ZrO2 is loose and porous, which is resulted from the volume expansion and phase transformation. B2O3 exhibits relatively low melting point (450 °C) and can form a continuous liquid layer on the sample surface, which would effectively inhibit oxygen diffusion. However, when the temperature increased above 1200 °C, B2O3 would be consumed completely owing to its severe evaporation and can no longer protect the matrix [7,8]. It has been reported that SiC introduction could improve the oxidation resistance of ZrB2 at high temperatures above 1200 °C by generating a borosilicate liquid layer [9-11]. ZrB2-SiC system has been extensively studied in the past and the optimum amount of SiC is confirmed to be 15-20 vol.% [12]. Nevertheless, it may not be the best choice to add SiC exclusively because of the following drawbacks: i. in the deep interior region of the bulk, where the oxygen partial pressure is low, SiC forms gaseous  products (e.g. SiO and CO) through active oxidation and thus a SiCdepleted layer is formed [13]; ii. the integrity of SiO2 liquid layer would be damaged when the temperature is above 1600 °C because of the transition from passive oxidation to active oxidation of SiC [14]. Metal silicides, such as MoSi2 and TaSi2, have also been introduced to ZrB2 to meliorate the oxidation resistance [15,16]. However, the melting points of metal silicides are relatively low, which limited their applications at high temperatures. In order to enhance the oxidation resistance of ZrB2, tungsten containing additives (e.g. WC, WB, WSi2 and W) were introduced. Zhang et al. [17-19] studied the oxidation resistance of ZrB2-WC ceramics and found that WC addition effectively reduced the weight gain and oxide samples at 1500-1600 °C. They confirmed that thickness of the presence of WO3 in the oxide scale resulted in liquid-phase sintering of ZrO2 and made the microstructure denser in comparison to the noncontaining WC samples. Recently, Zou et al. [22] evaluated the ablation resistance of ZrB2-SiC-WC ceramics through oxyacetylene flame at 2400 °C and found that the addition of WC could avoid the formation of the SiC-depleted layer. Table 1 is the summary of literatures that reported the effect of tungsten containing additives on the oxidation resistance of UHTCs. WB is characterized with high melting point (2670 °C) and could provide B2O3 during oxidation. Therefore, WB can better improve the self-healing ability of UHTCs compared to WC and W, yet few studies have discussed about the function of WB in the oxidation resistant property of ZrB2-SiC composites. What’s more, few reporters focused on coating materials. In order to ameliorate the  ⁎ Corresponding authors. E-mail addresses: yrniu@mail.sic.ac.cn (Y. Niu), xbzheng@mail.sic.ac.cn (X. Zheng).  https://doi.org/10.1016/j.corsci.2019.04.034 Received 8 June 2018; Received in revised form 23 April 2019; Accepted 25 April 2019  Available online 29 April 2019 0010-938X/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'C. Li, et al.  Table 1  Summary of  literatures on tungsten-containing ZrB2 based ceramics.  Corrosion Science 155 (2019) 155-163  Composition  ZrB2-WC-B4C ZrB2-WC-B4C ZrB2-WC-SiC-B4C ZrB2-WC-B4C ZrB2-MeSi2 (Me = Zr, Mo, Ta, W) ZrB2-W-B4C ZrB2-WSi2 ZrB2-SiC-WC  Material  bulk ceramic bulk ceramic bulk ceramic bulk ceramic bulk ceramic bulk ceramic bulk ceramic  Fbrication method  Test conditon  pressureless sintering pressureless sintering pressureless sintering hot pressing hot pressing hot pressing spark plasma sintering  oxidation, 1500, 1600 °C for 1, 2, 3 h oxidation, 1500, 1600 °C for 1, 2, 3 h oxidation, 1500, 1600 °C for 1, 2, 3 h oxidation, 1200, 1300, 1500, 1650 °C for 15 min heated from 800 to 1600 °C oxidation, 1650 °C for 15 min ablation, 2400 °C for 300 s  Ref.  [17] [18] [19] [15] [20] [21] [22]  Fig. 1. XRD patterns of as-received powders (a) and as-sprayed coatings (b).  intrinsic brittleness of ceramics, UHTCs could be used as coating materials to protect thermal structure components such as C/C and C/SiC composites [23-26]. Especially, vacuum plasma spray (VPS) technology is suitable for the preparation of UHTCs coatings for the following reasons: i. the temperature of the central plasma flame is above 10,000 °C, which is effective to melt materials with high melting points; ii. the process is done in an inert gas atmosphere, which could avoid the introduction of oxygen impurity to some content; iii. the high efficiency and automation of VPS could ensure high coating quality [16,27-29]. To the best of our knowledge, few scholars have studied the oxidation resistant behaviors of ZrB2-SiC-WB ternary composite coatings. In this work, ZrB2-SiC coatings with different contents of WB were fabricated by VPS method. The oxidation resistance of the composite coatings was evaluated in ambient air at 1500 °C for different time. The phase composition and microstructure changes of the coatings were characterized in detail. The function and influence mechanism of WB addition on the composition and microstructure changes of the oxide scale were focused on. It is expected that this work would shed some light on the role of WB in optimizing the oxidation resistant properties of UHTC composites.  2. Experimental procedure  The ZrB2-20 vol.%SiC powders containing different contents of WB (0 mol%, 5 mol%, 10 mol% and 15 mol%, noted as ZS, ZSW5, ZSW10 and ZSW15, respectively) were sprayed on graphite substrates via a vacuum plasma spray system (A-2000, Sulzer Metco AG, Switzerland) equipped with an F4-VB torch. Then the graphite substrates were removed by mechanical method to get free-standing coating samples for this research. The thickness of these obtained coatings was about 1.0 mm and the coatings were cut to an average size of 5.0 mm × 5.0 mm × 1.0 mm for later experiments. The coatings were placed in a tube-type furnace and were heated up to 1500 °C at a rate of 10 °C/min. The samples were incubated for 1 h, 3 h and 6 h respectively, and then cooled to room temperature naturally. The mass change of samples was measured by an electronic balance with an accuracy of 0.1 mg. The mass gain rate Wm% was  calculated as:  m  2  W %  m  =  −  m  1  m  1  100%  ×  the measured weight before and after the  where m1 and m2 represent oxidation, respectively. The phase compositions of the powders, as-sprayed and oxidized coatings were identified by X-ray diffraction (XRD, RAX-10, Rigaku, Japan) with Cu Kα (λ = 1.5406 Å) radiation. The microstructure and chemical composition of the as-sprayed and oxidized coatings were characterized by scanning electron microscopy (SEM, Magellan 400, FEI, UK) equipped with X-ray energy-dispersive spectroscopy (EDS, INCA SERIES, Oxford Instrument, UK). The porosity of the as-sprayed coatings were evaluated by three cross-sectional images with a magnification of 1000× using an image analysis software (Leica Qwin, Germany).  3. Results and discussion  3.1. As-sprayed coatings  The XRD patterns of the as-received powders and as-sprayed coatings are presented in Fig. 1. The peaks corresponding to WB were observed for the ZSW5, ZSW10 and ZSW15 powders. The major peaks of hexagonal ZrB2 were indexed in the ZS powders (PDF card number: 340423). Besides, it is worth mentioning that the ZrB2 peak of the WBshifted to a higher 2θ angle (Fig. 1a). containing composite powders The main reason is that the W substitution for Zr in ZrB2 reduced the average unit cell size, as the covalent radius of W (1.4 Å) is smaller than that of Zr (1.6 Å) [17,19]. These results indicated that a solid solution, (Zr, W)B2, was formed in the ZrB2-SiC-WB composite powders. Fig. 1b shows the XRD patterns of the as-sprayed coatings. The peaks corresponding to ZrB2 and ZrO phases were observed, while SiC and WB phases were not detected. Although they were not observed in the XRD patterns, the following microstructure and element mapping results have confirmed the existence of WB additive in the composite coatings. The surface and fracture morphologies of the as-sprayed coatings  156  \\x0c', 'C. Li, et al.  Corrosion Science 155 (2019) 155-163  Fig. 2. Surface morphologies of ZS, ZSW5, ZSW10 and ZSW15 coatings.  Fig. 3. Fracture morphologies of ZS, ZSW5, ZSW10 and ZSW15 coatings.  Fig. 4. Cross-sectional morphology and element maps of ZSW10 coating.  157  \\x0c', 'C. Li, et al.  Corrosion Science 155 (2019) 155-163  Fig. 5. Macroscopic photos of ZS, ZSW5, ZSW10 and ZSW15 coatings after oxidation at 1500 °C for 1, 3, 6 h.  were analyzed by SEM as shown in Figs. 2 and 3, respectively. All coatings were built up by melted and semi-melted particles. The porosity of the composite coating was calculated using the cross-section SEM images, which were about 15-20%. A cross-sectional microstructure image and element maps of the ZSW10 coating are revealed in Fig. 4. It can be seen that the elements of Zr, B, Si and W were uniformly distributed. Combined the XRD and EDS results, it is inferred that ZrB2, SiC, WB and (Zr, W)B2 existed in the ternary composite coatings.  3.2. Oxidized coatings  3.2.1. Macroscopic changes after oxidation tests The macroscopic morphologies of the samples after oxidation at 1500 °C for 1 h, 3 h and 6 h are shown in Fig. 5. The surface of samples changed from dark color to white color and the white region enlarged with the oxidation time increasing. The surface color of the ZS and ZSW15 coatings became white after 1 h oxidation, and the ZSW5 coating was completely changed to white color after 3 h oxidation. The ZSW10 coating kept the original dark color after 6 h oxidation. The mass gain rates of the composite coatings after oxidation for different time at 1500 °C are plotted as a histogram in Fig. 6. The mass gain rate of the ZS coating was about 20.9% after 1 h oxidation, which was far greater than other samples under the same oxidation  Fig. 6. Specific weight changes of ZS, ZSW5, ZSW10 and ZSW15 coatings at 1500 oC for 1, 3, 6 h.  158  conditions. And much smaller weight gain changes were observed in the WB-containing coatings. This suggested that WB was effective in improving the oxidation resistance of the ZrB2-SiC coating. It is worth noting that the mass gain rate of the ZSW10 coating was about 18.9% after 6 h oxidation, showing the lowest weight gain among all samples.  3.2.2. Oxidation at 1500 °C for 1 h The fractured SEM images of the oxide scales of the coatings after 1 h oxidation are showed in Fig. 7. For the ZS and ZSW15 coatings, no liquid layer was formed on the surface, while some liquid was observed inside the ZS coating (Fig. 7a). A liquid layer appeared on the surface of the ZSW5 and ZSW10 coatings. It can be seen that a thicker liquid layer could form in a short time for the ZSW10 coating (about 15 μm for the first one hour). The composition of the liquid layer was paid attention to as well. The liquid layer in the ZSW10 coating was mainly composed of SiO2, according to the EDS analysis (Fig. 7e). The scale under the liquid layer contained Zr, W, O, Si and B elements. It is worth noting that W and B mainly existed in the internal part owing to the evaporation of WO3 and B2O3.  3.2.3. Oxidation at 1500 °C for 3 h The phase compositions of the coating surface after 3 h oxidation were characterized by XRD, as shown in Fig. 8. It is discovered that the main composition was ZrO2. Within the resolution limit of XRD, no other crystalline phases were detected. The fracture morphologies of the ZS and ZSW10 coatings after 3 h oxidation are present in Fig. 9. A thin liquid layer and some needle-like particles were observed in the ZS coating (Fig. 9a). The needle-like particles (point A) was composed of ZrO2 while the liquid layer area (point B) was mainly composed of SiO2, as confirmed by the EDS results. The SiO2-rich layer was formed after 3 h oxidation in the ZS coating, indicating that the liquid content gradually increased with the oxidation time prolonging. A liquid layer with a thickness about 50 μm was formed in the ZSW10 coating, which was the thickest among all coatings (Fig. 9b).  3.2.4. Oxidation at 1500 °C for 6 h The fracture morphologies of the four kinds of coatings after 6 h oxidation are shown in Fig. 10. The liquid layer disappeared and some holes were observed inside the ZS coating. ZrO2 grains in columnar shape were observed in the holes from the magnified picture (Fig. 10a). A discontinuous liquid layer was formed on the surface of the ZSW5 coating (Fig. 10b). A continuous glass layer covered the surface of the ZSW10 coating, the thickness of which was about 40 μm (Fig. 10c). In contrast, the microstructure of the ZSW15 coating was very loose  \\x0c', 'C. Li, et al.  Corrosion Science 155 (2019) 155-163  Fig. 7. Fracture morphologies and element maps of the coatings after oxidation for 1 h: (a) ZS coating, (b) ZSW5 coating, (c) ZSW10 coating, (d) ZSW15 coating and (e) the element maps of ZSW10 coating.  (Fig. 10d). It is inferred that the ZSW15 coating was destroyed by bubbles due to the accumulation of gases. In addition, it is observed that some internal grains in darker color were wrapped by grains in whiter color in the ZSW5 coating (Fig. 10e). The EDS analysis indicated that the region B, C and D were composed of ZrO2, WO3 and borosilicate glass, respectively. This result suggested that ZrO2 was wrapped by WO3. Silvestroni et al. [15,30] also found that (Zr, W)B2 solid solution was formed in ZrB2-WSi2 composite ceramics and a coreshell structure, of which the core was ZrB2 and the shell was (Zr, W)B2, was detected by TEM. Therefore, the unique structure of WO3-wrapped ZrO2 was formed after oxidation for the ternary composite coatings.  3.3. Discussion  3.3.1. Thermodynamic analysis on function of WB and WO3 The oxidation process of ZrB2-SiC-WB composite coatings, which could take place in ambient air under high temperature (1500 °C), followed the reactions below [7,13,21,22]:  ZrB2(s) + 5/2O2(g) = ZrO2(s) + B2O3(l)  SiC(s) + O2(g) = SiO2(l) + C(s)  SiC(s) + 3/2O2(g) = SiO2(l) + CO(g)  WB(s) + 3/4O2(g) = W(s) + 1/2B2O3(l)  (1)  (2)  (3)  (4)  159  \\x0c', 'C. Li, et al.  Fig. 8. X-ray diffraction patterns of ZS, ZSW5, ZSW10 and ZSW15 coatings after oxidation at 1500 °C for 3 h.  WB(s) + 9/4O2(g) = WO3(s) + 1/2B2O3(l)  (5)  The oxidation processes and the oxidation products have great influence on the oxidation behavior of the composite coatings. We try to elaborate the effects of WB and its oxidation product (WO3) based on thermodynamic analysis in the following part.  3.3.1.1. Function of WB. Fig. 11 describes the oxygen partial pressure change as a function of temperature for the oxidation of ZrB2 and WB. The blue line indicates that ZrB2 is oxidized into ZrO2 and B2O3 (Reaction (1)) and the red line indicates the oxidation of WB into WO3 and B2O3 (Reaction (5)). It is demonstrated that the required oxygen partial pressure (pO2) for ZrB2 oxidation is lower than that of WB at 1000-2000 °C temperature range, namely, the oxidation of WB requires higher oxygen content than ZrB2 under the same condition. Therefore, WB is more stable than ZrB2 in an environment of low oxygen partial pressure. The volatility diagram of ZrB2-SiC-WB system at 1800 K was calculated based on NIST-JANAF thermochemical tables [31], as depicted in Fig. 12. It can be seen that the equilibrium oxygen partial pressures (pO2) for SiC oxidized to SiO2 (Reaction (2)) and for ZrB2 oxidized to ZrO2 (Reaction (1)) are 10−10.48 Pa and 10−10.38 Pa, respectively.  Corrosion Science 155 (2019) 155-163  When pO2 is below 10−8.59 Pa, WB is stable. When pO2 is between 108.59-10−4.67 Pa, W and B2O3 are produced (Reaction (4)). When pO2 is above 10−4.67, WB will be oxidized into WO3 and B2O3 (Reaction (5)). It can be seen that WB is the most stable substance at low oxygen partial pressure compared with ZrB2 and SiC.  3.3.1.2. Function of oxidation product WO3. The formation of the liquid layer on the coatings’ surface is vital for their oxidation resistance. The thickness of the liquid layer is one critical factor affecting the oxidation resistance of ZrB2 based composite coatings. The thickness of the liquid layer is related to the formation (generation) rate and consumption (vaporization) rate of borosilicate glass. If the formation rate is faster than the consumption rate, it means that a thicker liquid layer could form on the surface, which can more effectively hinder the transmission of oxygen to the coating interior. If the consumption rate is faster than the formation rate, it means that the liquid layer becomes thinner or broken, which could not effectively protect the matrix anymore. Compared with ZrB2-SiC coating, the addition of WB resulted in a thick liquid layer (Figs. 7, 9 and 10) and a reduction of weight gain (Fig. 6), for the reason that appropriate amount of WO3, the oxidation product of WB, could significantly reduce the volatilization rate of borosilicate glass. Cations are commonly regarded as network modifier in glass network. Cation field strength usually has direct impact on the arrangement of molecules for the liquid glass. The cation field strength is defined as Z/r2, where Z is the atomic number and r is the ionic radius (Å). The cation field strengths of groups IVB, VB and VIB elements are labeled in Fig. 13. It can be seen that V5+, Mo6+ and W6+ had relatively high field strengths of 23.62, 17.23 and 16.67 Å−2, respectively. While the cation field strengths of Zr4+ was the lowest among all elements (7.72 Å−2). The high cation field strength of W6+ would attract nonbridging oxygen around them, resulting in the formation of SiO2 phase separation and then the increasing of the liquid viscosity. Based on the Stoke-Einstein relation, the high viscosity glass could also help decrease the oxygen diffusion rate [32]. It has been reported that metal oxide additions could modify the [33-35]. Certain structure of B2O3, and then the stability of B2O3 amount of WO3 is beneficial for decreasing the evaporation rate of B2O3 and then improving the stability of the liquid layer. However, the existence of more WO3 would decrease the stability and enhance the evaporation of B2O3. Fig. 14 reveals the partial pressures of some oxidation products under different temperatures, including WO3, B2O3, SiO2 and ZrO2. It can be seen that the vapor pressures of B2O3 and WO3  Fig. 9. Fracture morphologies of ZS (a) and ZSW10 (b) coatings after 3 h oxidation.  160  \\x0c', 'C. Li, et al.  Corrosion Science 155 (2019) 155-163  Fig. 10. Fracture morphologies of coatings after 6 h oxidation: (a) ZS coating, (b) ZSW5 coating, (c) ZSW10 coating, (d) ZSW15 coating and (e) typical microstructure of ZSW5 coating.  Fig. 11. Thermodynamic oxygen pressures.  stability diagram of ZrB2 and WB under different  Fig. 12. Volatility diagram of ZrB2-SiC-WB system at 1800 K.  161  \\x0c', 'C. Li, et al.  Corrosion Science 155 (2019) 155-163  (Fig. 10a), due to the higher stability of WB compared with ZrB2 in the low oxygen partial pressure environment. Both the thick liquid layer and relative dense solid oxide layer contributed to the enhanced oxidation resistance of the ZSW5 and ZSW10 coating.  3.3.2.3. For ZSW15 coating. No liquid phase was observed on the surface of the ZSW15 coating after 1 h oxidation (Fig. 7d) and the oxide layer become more loose with many big holes after 6 h oxidation (Fig. 10d). It is inferred that the excessive WO3 formation will reduce the liquid stability and promote the liquid evaporation. What’s more, the accumulation of large amount of gases resulted in loose solid oxide layer. That is, the excessive amount of WB addition would lead to the degradation of the oxidation resistance.  4. Conclusions  ZrB2-SiC-WB ternary composite coatings were fabricated by vacuum plasma spray and their oxidation resistance and microstructure changes at 1500 °C were studied. Based on the experimental results and theoretical analyses, the following conclusions can be drawn:  (i) Appropriate amount of WB addition could effectively improve the oxidation resistance of ZrB2-SiC coating. While excessive of WB addition would reduce the oxidation resistance of ZrB2-SiC coating. The ZSW10 coating exhibited the lowest weight gain and thickest liquid layer among all samples. (ii) WB is a better boron source compared with ZrB2 for promoting self-healing ability. Based on thermodynamic analysis, WB is the most stable substance compared with ZrB2 and SiC in the low oxygen partial pressure environment. (iii) The formed WO3 oxidation product plays great function on the oxidation behavior of the ZrB2-SiC coating: the appropriate dissolution of WO3 could increase the viscosity of the liquid and thus decrease the oxygen diffusion; WO3 with optimum content could stabilize borosilicate glass and reduce its evaporation; the eutectic of WO3-ZrO2 phases contributed to the formation of a denser ZrO2 layer to impede oxygen diffusion.  Data availability  The research data supporting this publication are directly available within this publication.  Acknowledgements  This work was supported by the National Natural Science Foundation (for Young Scholar) of China under Grant 51102267, Engineering case study in extreme conditions using system mechanics approach (XDB22010202) and Youth Innovation Promotion Association CAS (2014223).  References  [2]  [1] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for extreme environments, Scr. Mater. 129 (2017) 94-99. F. Monteverde, L. Scatteia, Resistance to thermal shock and to oxidation of metal diborides-SiC ceramics for aerospace application, J. Am. Ceram. Soc. 90 (2007) 1130-1138. [3] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [4] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, High-strength zirconium diboride-based ceramics, J. Am. Ceram. Soc. 87 (2004) 1170-1172. [5] A.K. Kuriakose, J.L. 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According to ZrO2-WO3 phase diagram, WO3 could form a eutectic with 75 mol% ZrO2 above 1231 °C [36]. It is inferred that WO3 and ZrO2 could produce some liquid as well to fill holes and cracks, therefore a relatively dense scale could be formed to further impede oxygen diffusion (Fig. 7d).  3.3.2. Oxidation mechanism analysis of coatings 3.3.2.1. For ZS coating. Some liquid was formed after 1 h oxidation (Fig. 7a). Then a thin continuous liquid layer formed on the coating surface after 3 h oxidation (Fig. 9a), while the liquid phase disappeared and some holes were generated after 6 h oxidation (Fig. 10a). It can be concluded that the disappearance of liquid is the main reason for the poor oxidation resistance.  3.3.2.2. For ZSW5 and ZSW10 coatings. A continuous liquid layer was formed on the surface of both coatings after 1 h oxidation (Fig. 7b and c). 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Binner, Thermoablative resistance of ZrB2-SiC-WC ceramics at 2400 °C, Acta Mater. 133 (2017) 293-302.  [21]  [22]  [16]  [17]  [29]  [27]  [26]  [25]  [24]  [23]  C. Bartuli, T. Valente, M. Tului, Plasma spray deposition and high temperature characterization of ZrB2-SiC protective coatings, Surf. Coat. Technol. 155 (2002) 260-273. C.A.A. Cairo, M.L.A. Graca, C.R.M. Silva, J.C. Bressiani, Functionally gradient ceramic coating for carbon-carbon antioxidation protection, J. Eur. Ceram. Soc. 21 (2001) 325-329. E.L. Corral, L.S. Walker, Improved ablation resistance of C-C composites using zirconium diboride and boron carbide, J. Eur. Ceram. Soc. 30 (2010) 2357-2364. S. Tang, J. Deng, S. Wang, W. Liu, Comparison of thermal and ablation behaviors of C/SiC composites and C/ZrB2-SiC composites, Corros. Sci. 51 (2009) 54-61. Y. Niu, L. Huang, C. Zhai, Y. Zeng, X. Zheng, C. Ding, Microstructure and thermal stability of TaSi2 coating fabricated by vacuum plasma spray, Surf. Coat. Tech. 279 (2015) 1-8. [28] H. Pu, Y. Niu, C. Hu, G. Wang, H. Li, Y. Zeng, X. Zheng, Ablation of vacuum plasma sprayed TaC-based composite coatings, Ceram. Int. 41 (2015) 11387-11395. C. Zhai, Y. Niu, L. Huang, H. Pan, H. Li, X. Zheng, J. Sun, Microstructure characteristics and oxidation behavior of vacuum plasma sprayed tungsten disilicide Int. 42 (2016) 18798-18805. coating, Ceram. L. Silvestroni, D. Sciti, TEM analysis, mechanical characterization and oxidation resistance of a highly refractory ZrB2 composite, J. Alloys Compd. 602 (2014) 346-355. [31] M.W. Chase Jr., NIST-JANAF Thermochemical Tables, Institute of Physics, Woodbury, New York, 1998. [32] D.D. Jayaseelan, E. Zapata-Solvas, R.J. Chater, W.E. Lee, Structural and compositional analyses of oxidised layers of ZrB2-based UHTCs, J. Eur. Ceram. Soc. 35 (2015) 4059-4071. [33] R. Akagi, N. Ohtori, N. Umesaki, Raman spectra of K2O-B2O3 glasses and melts, J. Non-Cryst. Solids 293 (2001) 471-476. [34] D. Maniu, T. Iliescu, I. Ardelean, R. Ciceo-Lucacel, M. Bolboaca, W. Kiefer, Raman study of B2O3-SrO-CuO glasses, Vib. Spectrosc. 29 (2002) 241-244. [35] D. Maniu, T. Iliescu, I. Ardelean, S. Cinta-Pinzaru, N. Tarcea, W. Kiefer, Raman study on B2O3-CaO glasses, J. Mol. Struct. 651 (2003) 485-488. L.L.Y. Chang, M.G. Scroger, B. Phillips, Condensed phase relations in the systems ZrO2‐WO2‐WO3 and HfO2‐WO2‐WO3, J. Am. Ceram. Soc. 50 (1967) 211-215.  fourth ed., American  [36]  [30]  163  \\x0c']"
},{
  "_id": 59,
  "PDF": "Effects of heat treatments on oxidation resistance and mechanical properties of ultra high temperature ceramic coatings.pdf",
  "Text": "[\"Available online at www.sciencedirect.com  Surface & Coatings Technology 202 (2008) 4394 - 4398  www.elsevier.com/locate/surfcoat  Effects of heat  treatments on oxidation resistance and mechanical properties of ultra high temperature ceramic coatings  Mario Tului a ,⁎, Stefano Lionetti a , Giovanni Pulci b , Elviro Rocca b , Teodoro Valente b , Giuliano Marino c  a Centro Sviluppo Materiali S.p.A. — Via di Castel Romano 100, 00128 Rome, Italy b Rome University “la Sapienza”, ICMA Department — Via Eudossiana 18, 00184 Rome, c CIRA S.C.p.A. Via Maiorise, 81043 Capua,  Italy  Italy  Available online 8 April 2008  Abstract  A critical point  in the  development of  future  generation space vehicles  is  the  availability  of materials  able  to withstand the  extreme  temperatures generated during orbital  re-entry. UHTC (Ultra High Temperature Ceramics) materials, such as Zirconium diboride based ceramic  composites, exhibit outstanding oxidation and thermal shock resistance, high emissivity and very high melting temperature: all  these properties  make ZrB2 based materials possible candidates in designing thermal protection systems for the hottest structures of re-entry vehicles such as wing leading edges and nose-tip. An innovative, proprietary way to deposit ZrB2 based coatings by plasma spraying has been developed. In the present work, the influence of high temperature exposure on oxidation resistance and mechanical properties has been investigated. Several coatings  constituted by different amounts of SiC particles dispersed in a ZrB2 matrix were produced; some samples were tested in as sprayed conditions, other ones were submitted to a high temperature heat treatment before testing. Oxidation resistance was determined by exposing samples in air at  2073 K for 180, 1800 and 2520 s. Oxidation kinetics were analyzed and modelled. Evaluation of flexural strength and elastic modulus by means of  four-points bending tests were carried out on samples  showing better oxidation behaviour; mechanical  tests were performed also on samples  exposed at a typical operative temperature (1800 s at 2073 K). Results showed that heat  treatment does not affect oxidation resistance, whilst  it  significantly improves mechanical properties.  © 2008 Elsevier B.V. All  rights reserved.  Keywords: Ultra High Temperature Ceramics; Zirconium diboride-silicium carbide composites; Space vehicles materials; High pressure plasma spraying  1. Introduction  Current studies on re-entry technologies in aerospace industry are focused on Reusable Launch Vehicles (RLV) in order to reduce costs: this approach needs a technical effort to overcome more severe flight conditions [1-3]. Traditional materials can't withstand new mission profiles since they are subject to very high temperatures (up to 2300°K) in high oxidising environment; for this reason, it's necessary to develop new solutions in terms of materials or systems (bulk + coating) which have the required mechanical properties and offer an opportune protection to the flying hot structure [4,5]. Moreover for these RLVs it is very important to have thermal protection systems (TPS) that  ⁎ Corresponding author. Tel.: +39 06 50 55 742; fax: +39 06 5055 488.  E-mail address: m.tului@c-s-m.it (M. Tului).  0257-8972/$ see front matter © 2008 Elsevier B.V. All  rights reserved.  doi:10.1016/j.surfcoat.2008.04.015  are reusable too: it means that the TPS should be able to maintain its characteristics even after several missions. An innovative and proprietary technology [6] was developed, to obtain thick composite coatings constituted by a diboride as a matrix and a dispersion of carbide particles. Such coatings are currently developed to be used as TPS in future generation launchers. For this kind of materials reusability is  Table 1  Sample labels  Treatment  status  As sprayed  Heat  treated  Composition  60 vol.% ZrB2-40 vol.% SiC  60-40AS 60-40HT  80 vol.% ZrB2-20 vol.% SiC  80-20AS 80-20HT  \\x0c\", \"M. Tului et al.  / Surface & Coatings Technology 202 (2008) 4394-4398  4395  considering a generic re-entry mission, a RLV would be exposed at high temperature for a short time (e.g., a few minutes); tests carried out at 2073 K up to 2500 s can be considered, with some limitations, as a cumulative exposure to several re-entry missions of shorter duration.  2. Experimental  Two different compositions of coatings were produced, i.e., 60 vol.% ZrB2-40 vol.% SiC and 80 vol.% ZrB2-20 vol.% SiC. The coating deposition procedure consisted in the following steps: (i) agglomeration of commercial powders of ZrB2 and SiC, in the proper concentration, by spray drying; (ii) deposition by HPPS (High Pressure Plasma Spray) under inert atmosphere (Ar) on a flat graphite substrate; (iii) mechanical removal of the substrate. More details about powder preparation and plasma spray deposition can be found elsewhere [8-11]. Samples of both the compositions were heat treated under inert atmosphere (Ar) at a temperature of 2173 K for 7200 s. In the following, different samples will be identified with a label indicating composition and heat treatment status, as described in Table 1. Agglomerated powders before spraying, as well the surfaces of sprayed samples, were analysed by XRD (X-Ray Diffraction), in order to determine eventual phase transformations caused by spraying. Fracture surfaces of samples belonging to all the typologies were observed by SEM (Scanning Electron Microscopy), in order to determine the variation of microstructure induced by the heat treatment with respect to as sprayed samples. Samples belonging to all the typologies were exposed at high temperature (2073 K) in air. Three different experiments were carried out, on as sprayed samples, varying exposure time (180, 1800 or 2520 s). After exposure, the samples' surface was analysed by XRD. Then, the samples were cut, embedded in resin, lapped, and observed by OM (Optical Microscopy). Micrographs were analysed by an image analyser software to determine the thickness of the oxidised layer. Obtained data were interpolated using a parabolic equation, according to an oxidation kinetic model [12,13]. Such a methodology to estimate oxidation behaviour, instead of the weight variation determination, was  Fig. 1. XRD spectra of 60-40 powder, of 60-40AS sample and of 60-40HT  sample.  very important, it means they should have good properties even after several exposures at high temperature. Previous works showed that exposition at mission temperatures produced a partial sintering and the formation of a superficial oxide layer [7]. It is supposed that sintering has a positive effect on mechanical properties of the coating, while the oxide layer has a detrimental effect, because it is an area subject to crack initiation. Experimentally the measured effect is the combined one: until now the two phenomena couldn't be separated. In this work two different coating compositions, both belonging to the ZrB2-SiC family, have been heat treated both in inert gas and oxidative atmosphere, and then they've been characterised from a point of view of mechanical and oxidation resistance to verify separately the effect of each treatment on the coating properties. Results underline that the samples heat treated under inert gas, after further exposure in high oxidising environment, have “as is” better mechanical properties than the coating in the condition. Analysis of the oxide layer thickness on the heat treated coatings after an oxidation cycle showed a different behaviour of the two compositions: anyway, in both cases the applicability of these coatings was demonstrated as reusable, without loss of their thermal protection function. In fact,  Fig. 2. Fracture surfaces of a 80-20AS sample (left) and of 80-20HT sample (right) observed by SEM.  \\x0c\", '4396  M. Tului et al.  / Surface & Coatings Technology 202 (2008) 4394-4398  On the other hand, oxidation layer thickness measurements were carried out by several authors on bulk UHTC samples and that allowed to compare the obtained results with the bibliographic data [15-17]. Specimens for mechanical tests were cut by sparking erosion and mechanical finishing. Some specimens were oxidised by exposing them in air at 2073 K for 1800 s. Four-points bending tests, according to the ASTM C 1161-02c standard (45 × 4 × 3 mm), were carried out both on as prepared specimens and on oxidised ones. During the tests, strain gauges were used to determine MOR (Module of Rupture) and MOE (Modulus of Elasticity).  Fig. 3. XRD spectra of 60-40AS after exposure at 2073 K for 180, 1800, and  3. Results  2520 s.  chosen because of the difficulties to collect all the fragments of the glass phase formed on the surface of the samples during the exposition at high temperature in air; moreover, very often samples reacted with the ZrO2 brick which sustained them in the furnace. Similar problems were noted also by other authors [14].  Fig. 1 shows XRD spectra of the powder with composition 60 vol.% ZrB2-40 vol.% SiC, of a 60-40AS sample and of a 60-40HT sample. Similar results were obtained with powder and samples with composition 80 vol.% ZrB2-20 vol.% SiC. Fig. 2 shows the fracture surfaces of 80-20AS sample and of 80-20HT sample. Similar results were obtained with 60-40AS and 60-40HT samples. Fig. 3 shows XRD spectra of 60-40AS  Fig. 4. Cross section of 80-20AS sample oxidised at 2073 K in air for 1800 s.  \\x0c', 'M. Tului et al.  / Surface & Coatings Technology 202 (2008) 4394-4398  4397  Table 2  Table 3  Oxidation layer thickness determined by image analyses after different  time of  Four-point bending test results  exposure at 1800 K  Samples  Thickness in µm  60-40HT 80-20AS 80-20HT  180 s  110  123  130  1800 s  145  147  181  Oxidation  status  Non oxidised  Oxidised  2520 s  163  184  176  Sample  MOR (MPa)  MOE (GPa)  Average  Std. deviation  Average  Std. deviation  60-40AS 60-40HT 80-20AS 80-20HT 60-40AS 60-40HT 80-20AS 80-20HT  75  204  93  145  46  93  37  45  8.9  10.7  4.6  18.8  10.0  11.5  11.8  7.1  54  262  84  278  88  140  107  120  6.0  6.0  3.3  14.0  20.5  9.5  9.1  20.4  after exposure at 2073 K for 180 and 1800, respectively. Similar results were obtained on samples belonging to the other typologies (i.e., 60-40HT, 80-20AS, and 80-20HT). Fig. 4 shows a cross section of an 80-20AS sample oxidised at 2073 K in air for 1800 s. Different layers formed on the material when exposed in air at high temperature can be observed: a) thin layer of SiO2; b) layer formed by SiO2 and ZrO2; c) intermediate layer, depleted in SiC; d) bulk ceramic. In the following, layers a) and b) will be jointly identified as oxidation layer. Table 2 shows the oxidation layer thickness determined by image analyses. As an example, Fig. 5 shows the curve obtained by interpolating, with a parabolic equation, data reported in Table 2 referred to the sample 60-40HT. Finally, Table 3 reports the results of bending tests.  4. Discussion  Despite that SiC decomposes before melting, the coatings maintain the same phase composition observed in the starting powders. It could be explained by the fact that ZrB2 and SiC form a stable liquid at a temperature (about 2500 K) which is lower than SiC decomposition temperature (about 3000 K) [18]. High temperature heat treatment of coatings results in a sintering effect; the microstructure looses the typical morphology of thermal sprayed coatings and the lamellar structure of splats cannot be identified any longer; grains grow up to several microns. Such a microstructure modification produces different effects on oxidation resistance and mechanical properties of coatings, as shown by experimental results. High temperature exposure in oxidising conditions promotes the growth of a layer  Fig. 5. Oxidation layer thickness of 60-40 HT samples as a function of time of  exposure: parabolic interpolation.  constituted by ZrO2 and SiO2 on the material surface; increasing the exposure time, the layer thickness increases, following a parabolic relationship, and the SiO2 content increases too. Such a behaviour can be observed on all the examined samples, independently from the SiC content in the starting powders or from the coating status (i.e., as sprayed or heat treated). From the point of view of the desired application, oxidation resistance test results show that the thickness of the part of the coating affected by the exposition at high temperature (2073 K), for an exposition time up to 1 h, is in the order of 100-200 µm. Such values, similar to the ones measured by other authors on bulk [14-16], allow the use of UHTC coatings as UHTC samples TPS for re-entry space vehicles. Examining the results of mechanical tests carried out on not oxidised samples, reported in the first rows of Table 3, it can be observed that heat treated samples present dramatically higher values of fracture resistance (MOR) and of elastic modulus (MOE). Mechanical property values of oxidised samples, vice versa, are significantly lower than the corresponding values obtained on not oxidised samples, despite that high temperature exposure causes a microstructure variation similar to the one observed on heat treated coatings (60-40HT and 80-20HT). It can be explained considering that oxidation produces surface defects which can initiate cracks.  5. Conclusions  Several samples of ZrB2-SiC plasma sprayed coatings, which differed for composition and post treatment status (as sprayed or heat treated) were submitted to high temperature oxidising conditions and to mechanical tests. Results showed that SiC content variation within the range 20-40 vol.% and heat treatment do not affect oxidation resistance. Vice versa, mechanical properties are significantly increased by heat treatment. Oxidation test results evidence the high potential of these coatings used as thermal protection systems: even if with some limitations, the oxidative cycles performed in air using a furnace give an indication of coating performance respect to in air high temperature exposition. Evaluation of the behaviour of these coatings in a real mission environment, taking into account the interactive effects of other properties (like emissivity, thermal conductivity and ablation resistance), will be performed in the next future by testing the developed coatings in a plasma wind tunnel facility [19-21].  \\x0c', \"4398  M. Tului et al.  / Surface & Coatings Technology 202 (2008) 4394-4398  Acknowledgment  This work has been carried out in the frame of the ASA (Advanced Structural Assembly) project, funded by the Italian space agency (ASI — Agenzia Spaziale Italiana).  References  Thermal  Spray  2001, New Surfaces  for  a New Millennium,  28-30  Maggio 2001, Singapore, ASM Int. Publ., 2001, p. 259.  [10] C. Bartuli, T. Valente, M. Tului, Surface and Coating Technology 155  (2002) 260.  [11] M. Tului, G. Marino, T. Valente, High Temperature Characterization of  an UHTC Candidate Materials for RLV's, European Space Agency, 2003,  p. 161, Special Publication SP-521.  [12] M.M. Opeka,  I.G. Talmy, J.A. Zaykoski, Journal of Material Science 39  (2004) 5887.  [1] M.O. Wolfe,  Proceedings  3rd  Symposium on Naval  Struct. Mech,  [13] E. Opila, S. Levine,  J. Lorincz,  Journal of Material Science 39 (2004)  Pergamon Press, New York, 1984. [2] S. Drawin, Ann. Chim. (Paris) 17 (7-8) (1992) 455.  5969.  [14] S.R. Levine, E.J. Opila, et al., Journal of the European Ceramic Society 22  [3] Report Boeing C.A. for NASA Evaluation of Civil Hypersonic Transport  (14) (2002) 2757.  Reliability, Langley Research Center, 1990.  [15] M.M. Opeka, I.G. Talmy, et al., Journal of the European Ceramic Society  [4] E.V. Clougherty, E.T. Peters, D. Kalish, Mater. Processes 70s: 15th Nat.  19 (13) (October 1999) 2405.  SAMPE (Soc. Aerosp. Mater. Process Eng.) Symp. Exhib., 1969, p. 297.  [16] M. Gasch, D. Ellerby, et al., Journal of Materials Science 39 (2004) 5925.  [5] L. Kaufman, NSWC TR, 1986, p. 86.  [17] F. Monteverde, A. Bellosi, S. Guicciardi, Journal of the European Ceramic  [6] M. Tului, T. Valente; European Patent no. 1241278 (2002); USA Patent no.  Society 22 (3) (March 2002) 279.  6761937 (2002).  [18] S.S. Ordan ' jan, A.I. Dimitrev, E.S. Moroskina, Izv. Akad. Nauk. SSSR,  [7] M. Tului, G. Marino, T. Valente, Proceedings of the Second International  Neorg. Mater. 10 (25) (1989) 1752.  Meeting on Thermal Spraying (2006), vol. 201 (5), 2006, p. 2103.  [19] S. Caristia, F. De Filippis, A. Del Vecchio, E. Graps, Proceedings of 54th  [8] T. Valente, C. Bartuli, G. Visconti, M. Tului, in: C. Berndt (Ed.), Thermal  International Astronautical Congress, Bremen, 2003.  Spray. Surface Engineering via Applied Research, ASM International,  [20] S. Caristia, F. De Filippis, A. Del Vecchio, C. Purpura, 4th European  Materials Park, OH, 2000, p. 837.  Symposium Aerothermodynamics for Space Applications, Capua, 2001.  [9] C. Bartuli, T. Valente, M. Tului,  in: C.C. Berndt, K.A. Khor, E.F.  [21] G. Russo, F. De Filippis, S. Borrelli, M. Marini, S. Caristia, Advanced  Lugscheider  (Eds.),  Proc.  International  Thermal  Spray Conference:  Hypersonic Test Facilities, edited by AIAA, 2002, p. 313.  \\x0c\"]"
},{
  "_id": 60,
  "PDF": "Effects of LaB6 addition on arc-jet convectively heated SiC-containing ZrB2-based ultra-high temperature ceramics in high enthalpy supersonic airflows.pdf",
  "Text": "['Corrosion Science 75 (2013) 443-453  Contents lists available at SciVerse ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effects of LaB6 addition on arc-jet convectively heated SiC-containing ZrB2-based ultra-high temperature ceramics in high enthalpy supersonic airflows  Frédéric Monteverde a,⇑  , Davide Alfano b, Raffaele Savino c  a National Research Council of Italy, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, b Italian Aerospace Research Center, Via Maiorise, 81043 Capua (CE), Italy c Department of Aerospace Engineering, University of Naples ‘‘Federico II’’, P.le Tecchio 80, 80125 Naples,  Italy  Italy  a r t i c l e  i n f o  a b s t r a c t  The performances of a ZrB2-SiC-LaB6 ultra-high temperature ceramic (UHTC) was investigated in high enthalpy supersonic air flow. The UHTC material reached and maintained steady-state radiative surface temperature of 1973 K (monitored by pyrometer) for 5 min, and survived the arc-jet plasma exposure without any optical evidence of mechanical damage. The oxide scale covering externally the sample evolved into different textures. Transient thermal analysis processed via CFD provided a good agreement between the numerical steady-state temperature and the temperature distribution obtained experimentally. The contemporary addition of LaB6 and SiC to ZrB2 had a detrimental effect on the overall oxidation resistance.  Ó 2013 Elsevier Ltd. All rights reserved.  Article history:  Received 21 December 2012 Accepted 19 June 2013 Available online 28 June 2013  Keywords:  A. Ceramic matrix composites B. SEM B. XRD B. Modelling studies C. Oxidation C. Oxide coatings  1. Introduction  The design of hypersonic vehicles needs leading edge components with sharp profile to improve flight performances like drag reduction, easier manoeuvrability, or larger cross-range options during atmospheric exit and re-entry [1]. At some stage of an hypersonic (re-entry) flight, the aforementioned sharp components are subject of very stressful heat fluxes in corrosive plasmas from atmosphere, and temperatures exceeding 2300 K for several minutes may become for these components a projected service condition [2]. Such extreme conditions limit the spectrum of possible candidates that nowadays are intensively searched within the class of ultra-high temperature ceramics (UHTC). UHTC typically include non-oxides with melting/decomposition temperatures in excess of 3300 K. Diborides of the group IV transition metals like ZrB2 and HfB2 (3518 K and 3653 K of melting point, respectively) are currently the base systems taken in greater consideration to answer this technological challenge. Besides melting points greater than 3500 K [3], transition metal diborides have also thermal conductivity above 50 W/(m K) up to 1500 K [4], giving them a great potential advantage over more traditional materials to operate in extreme environments such as those experienced  ⇑ Corresponding author. Tel.: +39 0546 699758; fax: +39 0546 46381. E-mail address: frederic.monteverde@istec.cnr.it (F. Monteverde).  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.06.029  on wing leading edges and nose-cones of hypersonic vehicles, or in propulsion systems of rockets and missiles [5,6]. Massive single-phase UHTC for high temperature structural applications in severe environments are affected by poor oxidation/ablation resistance [7], as well as lacking damage tolerance [8]. A substantial research effort has been made along the past two decades to design and develop ultra-high temperature resistant composite systems. As far as transition metal diboride (MB2) - based ceramics are concerned, several additions and compositional refinements are being attempted [8], but outstanding achievements were actually reached for SiC particulate additions in the range of 10-30 vol% [7,9]. The inclusion of SiC was recognized to significantly improve thermo-mechanical stability and oxidation resistance of MB2 matrices at temperature up to 1800 K by forming ‘‘in situ’’ protective amorphous, fluid borosilicate coatings on the external surfaces directly facing the oxidizing environment at high temperature [10]. In more recent years, there has been a renewed interest in UHTC, particularly in the regard of methods for really pushing UHTC capabilities towards the ambitious goal of ultra-high temperature regimes, i.e. above 1800 K. In fact, for hypersonic space applications, the candidate materials would sustain service temperatures in excess of 2300 K in combination to fast flying air flows: in such operating conditions, silicate-based melts of any kind are removed rapidly, so that any oxidation protection afforded by fluid glassy oxide coatings is largely lost. Eakins and co-authors  \\x0c', '444  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  recently published a systematic review [10] on the on-going research attempts to enhance the resistance to oxidation of the MB2-SiC through the ‘‘in situ’’ formation and growth (at high temperature) of solid refractory oxides. Such new idea has involved the use of additives based on rare-earth elements like Ta [11-13] or La [2,14-17]. By using LaB6 [14-16], the improved oxidation resistance of ZrB2-SiC matrices has been attributed to the formation of a refractory oxidation solid product like lanthanum zirconate La2Zr2O7 with a (cubic) pyrochlore structure which, thermodynamically stable up to 2570 K (i.e. its melting point), was recognized capable of retarding the inward diffusion of oxygen. Lanthanum zirconate-based coatings are currently investigated as thermal barrier coating for the next generation of gas turbines to further enhance their efficiency, i.e. higher temperature capabilities and longer lifetimes [18]. The first cited example involving LaB6 as additive [14] is the study of Zhang who reported the description of a test up to 2673 K for 10 min: by means of an oxyacetylene torch the authors investigated the effects on the resistance to oxidation of 10 vol% LaB6 added to ZrB2-20 vol% SiC. The study right mentioned highlighted a significant improvement of the resistance to oxidation compared to an additive-free ZrB2-20 vol% SiC matrix, thanks to the growth of a compact La2Zr2O7 external oxide. Conversely, Ping and co-authors [15] tested the same nominal composition ZrB2-20%SiC-10%LaB6 in conditions of stagnant air at 2073 K for 60 min. They concluded that the overall effect of the LaB6 introduction was to significantly alter the ability of the external oxide scale to protect the underlying bulk by lowering its eutectic temperature and increasing the concentration of oxygen vacancies: the impact of these two interfering phenomena was detrimental, resulting in a reduced oxidation resistance compared to a LaB6-free formulation. Finally, Jayaseelan [16] processed SiC-reinforced ZrB2 materials in static air for 60 min at 1873 K and suggested that rare earth element additions may be useful approach to improving the oxidation resistance of UHTC at intermediate temperatures in hypersonic air. Taking into account this controversial context, the present article aimed at providing new and original elements of clarification regarding the effects on the resistance to oxidation coming from the addition of LaB6 particulate into a ZrB2-SiC matrix. Differently from studies designed for ground-based testing facilities such as ordinary oxidizing furnaces [15,16] or oxyacetylene torches [14,19], the present article dealt with UHTC in the ZrB2-SiC-LaB6 system convectively heated using an arc-jet facility, under conditions recalling those experienced during hypersonic re-entries. A very important peculiarity passed through the testing temperature range above 2073 K in presence of significant quantities of fast flying (fully) dissociated gases which approach the surfaces of the component: in such a way the conditions of an hypersonic flight were replicated more realistically. Properties and performances of the selected formulations in the ZrB2-SiC-LaB6 system were compared to those of other ZrB2-15 vol%SiC formulations already investigated as thermal protection (passive) system in very similar (arc-jet) testing conditions.  2. Materials and methods  2.1. Materials  A ZrB2-SiC-LaB6 composite (hereafter labeled ZSL10) was prepared from commercially available powders supplied by H.C. Starck: ZrB2 grade B, SiC grade BF-12, and LaB6 grade C. Typical particle size ranges of the as-purchased powders were the following: (1.5-3) lm for ZrB2, (0.5-1) lm for b-SiC, and (2-3) lm for LaB6. Raw powders in volumetric proportions 75%ZrB2 + 15%SiC + 10%LaB6 were mixed for 24 h in polyethylene bottle, using sil icon carbide media and absolute ethyl alcohol. The slurry was then dried through rotary evaporator and sieved (mesh opening 250 lm). The dried powder mixture was cold compacted (147 MPa of uniaxial applied pressure) into a ‘‘green’’ pellet (32 mm diameter) for hot-pressing. The ‘‘green’’ pellet was positioned inside a graphite die whose internal walls were lined with BN-spray coated graphitized 0.75 mm thick foil. Hot-pressing of the ‘‘green’’ pellet was carried out under partial vacuum (20- 100 Pa), heating rate 15-20 K/min up to 2200 K and 30 MPa of applied pressure. An isothermal hold at 2200 K was extended for 15 min increasing the applied pressure up to 40 MPa. The temperature was measured by means of an optical pyrometer focused on the outer wall of the graphite die. At the end of the hold, the applied pressure was released and the hot-press was let cooling to room temperature. After removal from the graphite die, the as-sintered pellet was about 15 mm thick.  2.2. Microstructure, mechanical and thermal properties  The bulk density (dS) of the as-sintered material was measured by Archimedes’ method, and compared to the theoretical value (dTH) calculated through the rule-of-mixture of the starting formulation. Density values equal to 6.09, 3.19 and 4.73 g/cm3 were used for ZrB2, SiC and LaB6, respectively, to calculate dTH. Crystalline phases were identified by X-ray diffraction (XRD, mod. D8 Advance, Cu Ka radiation, Bruker-Germany). A polished section (final finish 0.25 lm) of the hot-pressed material was prepared and observed by scanning electron microscopy (SEM, mod. S360 Leica Cambridge, UK) equipped with an energy dispersive analyzer (EDS, INCA Energy 300, Oxford Instruments, UK). The fracture toughness (KIc) was evaluated using 25 mm \\x02 2 mm \\x02 2.5 mm (length \\x02 width \\x02 thickness) chevron-notched beams (CNB) in flexure (4 bars) on a 4-pt alumina bending fixture with 20 mm and 10 mm as outer and inner span. The bars were fractured with a crosshead speed of 0.05 mm/min and KIc was calculated according to the ‘‘slice model’’ of Munz et al. [20]. Adopting the same fixture, flexural strength at room temperature and at 25 mm \\x02 1773 K in air was measured on chamfered bars 2.5 mm \\x02 2 mm (length \\x02 width \\x02 thickness, respectively), using the standards EN 843-1 (room temperature) and EN 820-1 (high temperature) as guideline. Ten and five bars, whose long edges were chamferered, were tested at room temperature and 1773 K (in air), respectively. Elastic modulus (E) and Poisson ratio (m) were measured using the resonance frequency in bending method on a 30 mm \\x02 8 mm \\x02 0.8 mm (length \\x02 width \\x02 thickness, respectively) plate using an universal frequency analyzer (mod. 4194A, Hewlett-Packard, Yokogama, Japan). Heat capacity (CP) and thermal diffusivity (DTH) were experimentally determined using a laser flash diffusivity apparatus (mod. LFA427, NETZSCH Gerätebau GmbH-Germany) in dynamic argon atmosphere between room temperature and 1773 K: both the surfaces of the test sample, a disc 12.7 mm in diameter and 2 mm in thickness, were painted with colloidal graphite. Thermal conductivity (KTH) was calculated according to the following expression  K TH ¼ dS \\x01 C P \\x01 DTH  The temperature dependence according to the relation  dS ðT Þ ¼ d0 ð1 þ aV \\x01 DT Þ\\x001  ð1Þ  of  dS was  taken into  account  ð2Þ  where aV is volume coefficient of thermal expansion, DT is the temperature difference with room temperature, d0 is the density value at room temperature. The thermal expansion vs. temperature was determined from room temperature to 1573 K by means of a dilatometer (mod. STA409, NETZSCH Gerätebau GmbH-Germany) in a  \\x0c', 'F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  445  stream of argon (5 K/min of heating rate), on a 25 mm \\x02 2.5 mm \\x02 2 mm (length \\x02 width \\x02 thickness, respectively) bar. The linear coefficient of thermal expansion aL was converted to (isotropic) volume coefficient aV = 3aL. All the specimens used for the thermo-mechanical characterization were machined through conventional mechanical tools.  2.3. Arc-jet tests: plasma wind tunnel ground facility and experimental set-up  The UHTC specimen for the arc-jet test was shaped as an hemisphere. It was first cut out by electrical discharge machining from the hot-pressed pellet, and then finished down to the desired dimensions through mechanical machining. The hemisphere had a nominal radius of curvature (R) of 5 mm. To in-depth evaluate the effects of the overall LaB6 content, two additional formulations (vol%), ZrB2 + 15SiC + 15LaB6 (ZSL15) and ZrB2 + 15SiC + 20LaB6 (ZSL20), were processed according to the procedures described in Section 2.1, and one hemispheric specimen was machined for material ZSL15 and ZSL20. The experiments on the hemispheres were carried out in high enthalpy supersonic air flows using an arc-jet plasma wind tunnel ground facility. In order to reproduce the air composition, molecular oxygen was mixed downstream to a primary nitrogen jet. The supersonic nozzle downstream the mixing chamber has an exit diameter of 22 mm. The nominal Mach number (M), determined by the area ratio of the nozzle, was 3. The hemisphere was held by an alumina tube at a distance of 10 mm from the nozzle exit (Fig. 1). The specific total enthalpy H0 at the nozzle exit was evaluated through an energy balance on the system torch-nozzle, measuring the energy supplied to the gas by the electric arc-heater and the energy lost through the water cooling system. During test, the surface temperature of the hemisphere was continuously measured (±10% uncertainty) by a two-colour pyrometer (Infratherm ISQ5, Impac Electronic GmbH, Germany) at an acquisition rate of 100 Hz. The two-color pyrometer exploited two overlapping infrared wavelength bands at 0.7-1.15 lm and 1-1.15 lm to measure temperature up to 3300 K. The measurement area of the two-color pyrometer was a round spot, 3 mm in diameter (see Fig. 1). The two-color mode was used to measure the surface temperature, whilst the single color mode allows to determine the (directional) spectral emittance ek once the temperature is known, under the ‘‘grey body’’ hypothesis. Due to the extremely high thermal loads upon the hemispheric model, surface chemical reactions like oxidation can be responsible for changes in ek. The two-color pyrometer overcomes this limit measuring the true temperature in the area of a round spot (3 mm diameter). In combination with the two-color pyrometer, the temperature distribution of the specimen (see Fig. 1) was also measured from a lateral view by means of an  infrared thermo-camera FLIR mod. P40. Since this instrument detects the spectral radiation coming from the sample along the infrared band wavelength of 8-12 lm, the actual surface temperature can be assessed as long as the average emittance in that range of wavelengths is known. Once the local temperature had been measured with the ratio pyrometer at the measurement spot, that value was input to determine the spectral emittance in the range of the IR thermo-camera, and finally the surface temperature distribution can be defined. Due to the hemispherical shape of the specimen only the portions of the surface almost orthogonal to the IR thermo-camera field-of-view were analyzed with the necessary accuracy. More details about the method used to evaluate ek is described elsewhere [21]. The arc-jet test on the ZSL10 hemisphere (hereafter labelled T1) was conducted in order to reach and stabilize for a duration of about 5 min an equilibrium (surface) temperature around 1973 K controlled by the two-color pyrometer. The gas flow rate was 1 g/s, and the static chamber pressure about 200 Pa. Once ignited, the arc-jet was steadied at the nominal supersonic flow conditions, then the specific total enthalpy (H0) was set at 11.4 MJ/kg to reach the desired target temperature of 1973 K. During the 5 min hold, the maximum stagnation point pressure (PMAX) was about 8 kPa. Also the ZSL15 and ZSL20 hemispheres were tested under the same conditions in terms of gas flow rate, H0, PMAX and duration, with the aim of assessing the effects of increasing content of LaB6 (at partial expenses of lower total content of ZrB2). After arc-jet test T1, the oxidized hemisphere (external surface and polished cross section) was analysed by XRD and SEM-EDS. External surface profile and mass of the hemisphere were measured on the (un-treated) sample before arc-jet testing by using, respectively, a 2-lm radius diamond tip contact stylus Taylor Hobson (mod. Talysurf Plus) profilometer, and an analytical balance (10\\x002 mg of accuracy) and compared to values measured after arc-jet testing. An estimate of the radius of curvature (R*) was extrapolated from the best-fitting of the real curved surface, under the hypothesis of an ideally perfect round profile.  3. Theory of the thermal model and Computational Fluid Dynamics (CFD) simulations  While the surface temperature of the hemisphere (according to the experimental set-up described in Section 2.3) and stagnation pressure were directly measured, the composition of the reacting gas mixture at the sample’s surface was computed through a numerical model based on the solution of the Favre-averaged Navier-Stokes equations, taking into account gas species in thermochemical non-equilibrium. The CFD simulations were carried out using methods documented in Refs. [22,23], in which thermal and chemical non-equilibrium, surface catalysis properties,  (a)  (b)  t  i  x  e  e  l  z z  o  N  UHTC  hemisphere (1) Alumina holder (2)  Fig. 1. (a) Side-view of the UHTC hemisphere mounted on the alumina holder before testing; (b) the measurement spot of the two-color pyrometer (filled circle) focused on the hemispheric surface during arc-jet experiment.        \\x0c', '446  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  changes of the surface temperature during the test and the radiation emitted from the surface were all taken into account. The adopted energy balance model schematically requires a ‘‘local’’ surface energy balance at steady state, with the total aero-thermal heating of the surface (convective heat flux from the surrounding gas to the wall minus conductive heat flux through the material) equals the energy transferred away from the surface by radiation. Such surface energy balance model was adopted to compute the surface temperature and the hot-wall heat flux, given the conditions of flow, environment, and thermo-physical properties of the sample. The convective heat flux from the gas to the wall is given by a ‘‘Fourier’’ contribution related to the temperature gradient, and a ‘‘chemical’’ contribution related to the gradient of the concentrations of dissociated chemical species. The thermal model was based on the solution of the unsteady energy equation in the solid, with the surface heat flux updated at each iteration to account for the energy re-emitted radiatively and for the changes in convective heat flux due to changes in surface temperature. CFD simulations were applied on the only ZSL10 material.  4. Results  4.1. Microstructure and basic properties of the as-sintered materials  Once dried and sieved, the powder mixtures ZSL10, ZSL15 and ZSL20 were investigated by XRD and SEM-EDS. On one hand, XRD analyses did not detect, besides the starting compounds, any new extra (crystalline) phases. In addition, the broadening (i.e. the full width at half maximum) of the XRD reflections belonging to ZrB2 phase before and after wet ball-mixing did not change appreciably. It was inferred that the particle size range of ZrB2 did not vary significantly in consequence of the ball-mixing, drying and final sieving. On the other hand, a limited uptake of oxygen in the form of amorphous boron oxide and zirconia very likely took place due to the well-known consequence of mixing (even in anhydrous wet ambient) abrasive powders using hard mixing media. Ortiz in fact reported that X-ray photoelectron spectroscopy unambiguously detected the presence of zirconia and boron oxide in the as-purchased ZrB2 powder as well as after wet ball-mixing, more abundant in the second one [24]: specific SEM-EDS analyses in the present work confirmed this trend. Even though the three powder mixtures picked up an additional elemental amount of oxygen below 1 wt%, such contamination had no adverse consequence on the overall sinter-ability of the final powder mixtures. In fact, the selected hot-pressing conditions allowed to obtain fully dense LaB6-containing ZrB2-15%vol SiC samples with bulk density (dS) very close to the expected theoretical value (Table 1). The XRD analysis of the hot-pressed materials identified only ZrB2, LaB6 and a-SiC as crystalline phases. Compared to the starting b-SiC powder, the high temperature treatment activated the b ? a polymorphic transformation. Polished areas of the internal volume analysed via SEM did not reveal apparent residual porosity, but rather a homogeneous spatial distribution of the three phase ZrB2, LaB6, and SiC constituting the ceramic bulk (Fig. 2). Representative grain size ranges of 4-5, 2-2.5 and 2-5 lm were estimated for the phases ZrB2, SiC and LaB6, respectively. Other properties of the ZSL10 material were measured, summarized in Table 2 and com Fig. 2. SEM micrograph of a representative polished area of the as-sintered ZSL10 material: the main phases ZrB2 (Z, lighter grey level), SiC (S, dark spot) and LaB6 (L, darker grey level) are indicated.  pared to a ZrB2-15%SiC formulation already employed and tested as thermal protection system [25]. Flexural strength in static air at 1773 K for instance increased, whilst fracture toughness remained rather poor. Interestingly, values of thermal conductivity ranging from 129 W/(m K) (at 300 K) to 98 W/(m K) (at 1773 K) were measured: an outstanding thermal conductivity is desired feature for application in aerospace like wing leading edge for hypersonic vehicles. The hot-pressing conditions allowed also to fully densify both the formulations ZSL15 and ZSL20 (see Table 1). Inspections of the internal bulk by means of SEM (polished sections) counted negligible levels of residual porosity. No further properties were measured on the as-sintered ZSL15 and ZSL20 materials.  4.2. Arc-jet testing T1 on the ZSL10 hemisphere  The ZSL10 hemisphere was object of the arc-jet test labeled T1. The most important experimental data, including duration, specific total enthalpy (H0), measured temperature (TPYR) and directional spectral emittance ek are summarized in Fig. 3. During the hold, the variation of TPYR was almost negligible (within the instrument’s accuracy), indicating that the ceramic hemisphere reached thermal equilibrium. The values of ek slightly decreased vs time: this can be very likely interpreted as consequence of a gradual change in the chemical composition of the exposed surface. The ZSL10 hemisphere survived the arc-jet plasma exposure without any optical evidence of mechanical damage. At the same time, the interaction with the high-enthalpy supersonic air-flow gave rise to surface modifications. After arc-jet test T1, the ZSL10 hemisphere had a net specific mass change of (1.27 ± 0.01) mg/ cm2, assumed that all the modifications entirely occurred upon the round surface of the sample: this corresponds to a net mass gain of 0.3%. The radius of curvature R*, estimated before and after arc-jet test T1, was averaged along two orthogonal linear scans of the initial R* (i.e. RI⁄) the entire curved surface: changed from 5.020 ± 0.005 mm to 5.14 ± 0.01 mm, so that, based on these values, bulk recession was inferred not taking place. The modified R* after arc-jet test T1 (i.e. RF* = 5.14 mm) is direct consequence of  Table 1 Theoretical (dTH), bulk (dS) and relative (rd) density, residual open porosity (P) and typical grain size range (Dx).  Sample  ZSL10 ZSL15 ZSL20  dTH (g/cm3)  dS ± 0.02 (g/cm3)  5.50 5.42 5.36  5.48 5.41 5.36  rd (%)  99.7 99.8 100  Pa (%)  <0.1 <0.1 <0.1  Dx (lm)  ZrB2 4-5  a By SEM on polished section.  SiC 2-2.5  LaB6 2-5  \\x0c', 'Table 2 Elastic modulus (E), Poisson ratio (m, fracture toughness (KIc), flexural strength (r),  linear coefficient of thermal expansion (k) and thermal conductivity (KTH).  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  447  k (ppm/K)  20-1300 °C  6.95 6.68  KTH (W/(m K))  300 K  129 62.5  1073 K  106.7 65.6  1773 K  97.8 66.2a  E (GPa)  m  p  KIc (MPa  m)  r20°C (MPa)  r1500°C (MPa)  ZSL10 ZSM [25]  465 480  a Measured at 1473 K.  0.123 0.13  2.84 ± 0.4 4.07 ± 0.03  608 ± 103 887 ± 125  405 ± 30 255 ± 25  not regularly compact due to the presence of internal defects (holes, cavities, pits), and essentially is composed of two distinct sub-regions. In Fig. 5a (area-1), a first (outer) sub-region consists predominantly of a grained oxide with isolated spots of residual glass. SEM-EDS analyses confirmed that the (grained) oxide is zirconia, while the glassy product is a borosilicate (BSG) containing variable amounts of Si, O, La, B and Zr. In fact, it was further verified that the chemical composition of BSG changes, depending both on its distance from the exterior and on the relative position with respect to the symmetry axis. Again, a second (inner) sub-region shown in Fig. 5a, more porous than the outer one, underlines the un-reacted bulk, is made of only zirconia grains and remained depleted of glassy residues. In Fig. 5band c (area 2 and 3), morphology, size and distribution of the residual porosity inside the oxide scale changed due to differentiated equilibrium temperature. Such regions of the external surface (and of the inner volume) reached lower equilibrium temperature (see Fig. 4), compared to zones facing the hot plasma and closer to the stagnation point (i.e. area 1 in Fig. 5a). This permitted the resident BSG to better resist against boiling/volatilization, and thus to remain inside the oxide scale in larger amounts: an example of such configuration is shown in Fig. 6. In those regions like area 1 near the stagnation point (see Fig. 5a), the BSG ‘‘in situ’’ forming at the oxide scale/bulk interface, which is exposed to higher temperature compared to regions like area 2 and 3, suddenly begins to boil and volatize: volatile by-products internally produced search to escape to the exterior, leading to the formation of large internal pits and voids. Some phenomena of glass separation took place (Fig. 7). Independently from the location, and in contrast to other authors [2,14-16], by using XRD and SEM-EDS the formation of solid crystalline Lanthanum zirconate La2Zr2O7 has never been ascertained anywhere in the new-formed external oxide scale. Finally, regardless of the position in respect to the symmetry axis, no glassy coating was seen atop of the new-formed oxide.  4.4. Computational fluid dynamics simulations and thermal analysis results  During the isothermal hold of test T1 (300 s at 11.4 MJ/kg of H0), the directional spectral emittance at 1 lm (e1lm) mildly decreased  Fig. 3. Arc jet test T1: surface temperature TPYR (by two-color pyrometer) of ZSL10 hemisphere vs. time (t); directional spectral emittance at 1 lm (e1lm) and specific total enthalpy H0 are also indicated.  the growth of a new external oxide scale whose thickness, microstructure and basic chemical composition were the focus of SEM-EDS analyses in Section 4.3.  4.3. Microstructure modifications upon oxidation  The main consequence of the arc-jet test T1 was an obvious modification of the external surface which got an overall whitish coloration, while appearing still compact and rather smooth. The XRD analyses upon the (oxidized) external surface provided evidence of the presence of monoclinic zirconia (m-ZrO2) as crystalline phase. The overall prevalence of m-ZrO2 implies almost complete transformation from the tetragonal structure upon rapid cooling, which involves a volume expansion of about 3%. Actually, as a result of the uneven distribution of the equilibrium temperature upon the surface during the hold of test (Fig. 4), the microstructure of the oxide scale covering externally the sample evolved into different textures. Fig. 5 shows the half full cross-section of the ZSL10 hemisphere after test T1, in addition to some enlarged views of representative spots from the external oxide scale. The oxide scale, whose thickness ranges approximately from 180 lm (at the symmetry axis of the specimen) to 150 lm, is  T1  (a)  (b)  Experimental points by IR  thermo-camera +  two -color  pyrometer   ---Calculated by CFD  1700  1820  1940  2060  2180  2300  Fig. 4. Arc-jet test T1 (11.4 MJ/kg of specific total enthalpy): (a) equilibrium temperature T vs x coordinate upon the curved surface (calculated by CFD); (b) distribution map of T(K) throughout the bulk (by CFD).  ♦ ♦ \\x0c', '448  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  (a)  (a) oxide scale: 180 μμm  (c) oxide scale: 150 μm  (b)  25 μm  (c)  Fig. 5. ZSL10 hemisphere after arc-jet T1: SEM micrograph patchwork of the half-full oxidized polished cross-section and some enlarged views from area 1, 2 and 3; the magnification marker is applicable for (a-c) views.  (a)  (b) 1: Zirconia phase  Zr  O  Zr  2  1  O  (c) 2: BSG  20 μm  B  La  Al  Si  Zr  Fig. 6. (a) SE-SEM micrograph from the oxide scale/internal bulk interface of ZSL10 (see area 2 in Fig. 5) after arc-jet test T1; EDS spectra from point 1 (b) and 2 (c, BSG borosilicate glass).  from 0.6 to 0.54 (Fig. 3). The spectral emittance (ek) is coined as a term to describe the effective spectral emissivity for a real surface which, in case of rougher surface for instance, increases propensity for heat absorption. This range of values is slightly lower than the value of 0.66 measured for a ZrB2-15%SiC formulation tested in very similar experimental conditions [21]. The surface temperature over the hemisphere measured during the 5 min hold (infrared thermo-camera combined to pyrometer measurements), and the corresponding distribution calculated through CFD simulations are already reported in Fig. 4. Measurements in the portions close to the stagnation points were limited by the field-of-view: this explains why in Fig. 4 the experimental points are missing between 0 (front portion) and 1 mm as x coor dinate. In specific, e1lm in the range of 0.56-0.6 for temperature between 1900 and 1970 K (recorded by the two-color pyrometer) are plausible values when the outer layer is predominantly made of zirconia. Literature data reported 0.66 ± 0.03 as value of total hemispherical emittance at a temperature of 1865 ± 20 K [26]. The presence of an external oxide was taken into account to likely match the thermal model to the real final configuration of an ‘‘oxidized’’ hemispheric sample (see Fig. 5). The thermal model was schematized adding a continuous external scale upon the native round profile of the hemisphere, in accordance with the microstructural features described in Section 4.3. The presence of defects and residual BSG, which occupy very limited volumes compared to the entire bulk of the hemispheric sample, was not included into  \\x0c', 'F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  ZrB2 þ 2:5O2 ðgÞ ¼ ZrO2 ðsÞ þ B2O3 ðlÞ  2LaB6 þ 10:5O2 ðgÞ ¼ La2O3 þ 6B2O3 ðlÞ  La2O3 þ 2ZrO2 ¼ La2 Zr2O7  449  ð3Þ  ð4Þ  ð5Þ  The scheme behind such idea was to favor, during the exposure to arc-jet supersonic air flow, the reaction between La2O3 (i.e. the expected primary oxidation by-product of LaB6) with zirconia (i.e. the principal expected solid oxidation by-product of ZrB2) in view to form, according to the reaction (5), an solid scale predominantly made of La2Zr2O7 cubic pyrochlore (hereafter labeled LZ). The occurrence of reaction (5) would therefore to succeed in growing ‘‘in situ’’ a barrier against oxidation during high temperature treatment in oxidizing environment. In fact, the family of ceramic materials typified by the composition A2B2O7 was already identified having great utility as thermal barrier coatings on metallic substrates [18]. Pyrochlore materials have also phase relationship in which the pyrochlore structure is phase stable up to the melting point, about 2570 K in the case of La2Zr2O7. Based on XRD and SEM-EDS analyses conducted after arc-jet testing on the three formulations ZSL10, ZSL15 and ZSL20, the ‘‘in situ’’ formation of such self-generating protective oxidation barrier has not ended up successfully, and zirconia and a (La + Zr)-containing borosilicate glass (BSG) remained, upon cooling to room temperature, the almost exclusively present solid oxidation by-products. Some Al from the sample holder was incorporated into the BGS during high temperature exposure (see Fig. 6c). Likewise, the introduction of LaB6 was also thought to control the viscosity of BSG at high temperature by favoring the onset of immiscible liquids inside the glassy melts. Phase separation phenomena with flower-like or round patterns made of silica, as well as crystallized Lanthanum silicates, were only occasionally observed in the oxide scale of the ZSL10 sample (Fig. 7), more frequently increasing the content of LaB6 (i.e. sample ZSL15 and ZSL20, see Fig. 8) at partial expenses of less ZrB2. An extension of research activity was further undertaken on the selected ZSL10 formulation to discern if the formation of the LZ phase did not initiate because of special conditions occurring during arc-jet testing like very rapid temperature transients above 1900 K in severely reactive air flows for short exposure times. It is worth noting that the key difference for the use of materials in hypersonic flights is the relatively short exposure to very high temperature compared with thousands of hours required for protective coatings inside gas turbine engines. Internal research activities, not the object of this work, agreed with the conclusions of some authors [29,30] that, based on reaction (5), a mixture of zirconia and lanthana powders in molar ratio 2:1 is fully converted into LZ by heating in air at atmospheric pressure up to 1673 K for 60 min (300 K/h heating rate and rapid cooling).  Fig. 7. Glass separation (GS) and crystallization (CR) after arc-jet test T1 (SE-SEM micrograph).  in the oxide scale of ZSL10  the refinement of the thermal model. A global emittance of 0.6 for zirconia [27] was used: according to Marschall [28], although surface roughness may cause deviations from the predictions, the measured average e1lm is still the best approximation currently available for the global emittance of the UHTC specimen during testing. By using this updated thermal model, a transient thermal analysis processed via CFD in conditions of H0 = 11.4 MJ/kg provided a good agreement between the numerical steady-state temperature and the temperature distribution obtained experimentally through the two-color pyrometer and the IR thermocamera. Based on the numerical simulations, radiative steady-state surface temperatures decreasing from 2300 K to 1850 K were predicted (see Fig. 4a). In addition, by using CFD simulations, a raise of the shear stress up to 0.55 kPa for an x coordinate of 1 mm has also been calculated, with a further decrease down to 0.2 kPa. Due to the combined effect of shear stress and ultra-high temperature, the absence of residual BSG atop of the exposed round surface can be regarded as an expected result. An increased flow ability of a (less viscous) BSG due to the presence of La, Zr and B compared to pure silica glass made easier its removal from the external surface of the specimen at elevated temperatures.  5. Discussion  5.1. Thermodynamics vs. kinetic factors competing in the formation of La2Zr2O7 pyrochlore  Lanthanum hexaboride powder (LaB6) was intentionally added to ZrB2-15%SiC to take advantage from the following chemical reactions  Fig. 8. Glass separation and crystallization near the oxide scale/un-oxidized bulk interface (dashed line) of ZSL15 (a) and ZSL20 (b) after arc-jet testing (SE-SEM micrographs).  \\x0c', '450  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  ent arc-jet experiments, the expected formation of LZ did not take place, very likely for kinetics issues related to rapid transient heating/cooling conditions, as well as to relatively short hold at high temperature. It seems also very likely consider that lanthana coming from reaction (3) is not straightway available for reaction (5) but tends to be rapidly incorporated by the ‘‘in situ’’ forming BSG: this plays for reaction (5) as unfavorable condition to proceed forward. A continuing research activity is underway on revised ZrB2-SiC-LaB6 formulations to firmly assess that LZ phase during arc-jet test has not formed due to kinetics issues, i.e. related to rapid transient heating/cooling conditions, as well as to relatively short hold at high temperature.  Fig. 9. Multiple thermal treatments in ambient static air using the bottom-up (8 \\x02 8 \\x02 0.8) mm3 loading furnace and three plates of material ZSL10 (O-C: opening-closure of the chamber to insert/take out samples).  5.2. Was the presence of LaB6 in the ZrB2-SiC base system effective against oxidation?  The focus of the slowest mechanism governing oxidation in conditions typical of an arc-jet experiment was in-depth reasoned. Very recently, Parthasarathy modeled oxidation kinetics of SiCcontaining refractory diborides based on experimental data reported on literature in the temperature regime of 1500-2500 K [31]. Parthasarathy in Ref. [31] concluded that data obtained using arc-jet tests fall well outside the model predictions, indication that some aspects of arc-jet conditions were not fully captured in the adopted model. In the present study, a number of fundamental evidences like (i) the absence of a glassy film coherently covering the external surface of the sample, (ii) the absence of an evident internal SiC-depleted region, (iii) the presence of several defects (i.e. pits, voids) in the oxide scale, and (iv) the internal formation of a BSG not uniformly distributed and much more fluid (compared to pure silica) led the authors of the present work to indicate the diffusion of oxygen through the zirconia skeleton (that constitutes the supporting frame of the external oxide scale) as the slowest mechanism controlling the oxidation rate of the ZrB2-SiC-LaB6  system in such special environment. x  The increasing addition of LaB6 (x = 10, 15, 20) into a (100\\x00x)vol% ZrB2-15vol% SiC mixture had an increasingly detrimental effect on the overall oxidation rate by facilitating the inward diffusion of oxygen thanks to an enhanced concentration of oxygen vacancies in the zirconia-based oxide scale. At reduced oxygen partial pressure (condition that exists close to the interface between the external oxide scale and the underlying un-reacted bulk) zirconia becomes non-stoichiometric by forming oxygen lattice vacancies. Oxygen vacancies allow oxygen ion transport through the scale, faster than in presence of a fully stoichiometric zirconia, and thus enhance oxidation of the underlying bulk. Also, according to reaction (4), the copious injection of Boron (provided by LaB6) through the yield of B2O3 (l) into the ‘‘in situ’’ forming BSG greatly reduced its viscosity at high temperature: it follows that the internal gaps/channels separating zirconia grains, which are (partially) filled of the glassy medium, become more permeable to inward diffusion of oxidants. Improved oxidation resistance was sometimes related to the presence of transition metal oxides in the BSG which, inducing phase separation (immiscibility), may lead to increased liquidus temperatures and viscosities, i.e. reduced oxygen diffusion rate [10,32]. The characteristic feature of an immiscible liquid glass can thus be beneficial for retarding oxygen transport and suppressing (partially) the evaporation of the boria component from the ‘‘in situ’’ formed BSG. Although it was clearly observed that the increase of the LaB6/ZrB2 volume ratio boosted phase separation, this phenomenon played a role of marginal importance, and did not contribute to an effective slow-down of the diffusion rate of the oxidants through the oxide scale. Actually, our microstructural observations indicated that, under high enthalpy supersonic and highly dissociated air flows (which induced surface temperatures  Fig. 10. XRD patterns from plate 1, 2 and 3 of material ZSL10 after oxidation, according to thermal treatments in Fig. 9: only the most intense peaks for La2Zr2O7 pyrochlore (LZ), tetragonal (TZ) and monoclinic (MZ) zirconia are indicated. The dominant phase is MZ.  It was matter of reasoning in fact whether kinetic aspects versus favorable thermodynamic conditions dominated the competition of the LZ phase onset during the arc-jet experiments herein presented. For the present study, three plates 8 mm \\x02 8 mm \\x02 0.8 mm (side \\x02 side \\x02 thickness, respectively), were cut from ZSL10 material, cleaned, placed on zirconia spacers so as to have minimal mutual contact area and were then heated up to 1973 K in static ambient air using a bottom-up loading furnace according to the schedule depicted in Fig. 9. This type of multiple thermal treatment for the three ZSL10 plates was driven by the need of understanding if kinetic aspects more than thermodynamic factors prevented the formation of LZ phase. Regardless of the thermal inputs (see Fig. 9), all the three plates resulted heavily oxidized. Based on a visual inspection inside the plate-2 fractured manually, only a central inner core remained un-oxidized, while plate-1 and plate-3 have the internal bulk completely oxidized. XRD analyses on the external oxidized surfaces (Fig. 10) verified the formation of minor amounts of LZ only in the plate-3. The results of this experiment agree with the outputs of Jayaseelan [16] and, at the same time, do not contradict those emerging from the present arc-jet test. In effect, compared to the conditions of the arc-jet test in terms of duration and heating/cooling rate, much longer durations of (deliberately) slower thermal transients (see Fig. 9) resulted essential to initiate the formation of LZ. In the pres \\x0c', 'F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  451  approaching 2300 K), the contemporary presence of SiC and LaB6 additives does not bring obvious advantages to improve oxidation resistance. In fact, just after having formed BSG, a release of gaseous by-products concomitantly starts taking place: such gases search escape pathways to the exterior creating new accesses from/to the exterior, and contributed to deteriorate the overall physical integrity of the oxide scale. The creation of defects of several hundreds of microns at the oxide scale/un-oxidized bulk interface, or inside the oxide scale itself, was attributed to the abundant release of gases evolving out from BSG ‘‘pools’’ (Fig. 11). The term ‘‘pool’’ was coined as such only for sake of clarity to indicate a posttest feature of BSG (easily observable by SEM-EDS in those locations at the oxide scale/un-oxidized bulk interface) which very likely did not exceed its boiling point. Several incipient micro-bubbles are clearly visible inside the BSG ‘‘pools’’ (see Fig. 11). The adherence of the oxide scale to the underlying diboridebased bulk, chiefly for those zones depleted of glassy interlayers (i.e. BSG ‘‘pool’’), was adversely affected by the tetragonal-tomonoclinic transformation of the zirconia phase. In effect, at high temperature zirconia is tetragonal [33]. Upon rapid cooling to room temperature within few minutes tetragonal zirconia, in absence of ion stabilizers, transforms to monoclinic with volume expansion of 3%. Our SEM-EDS analyses, according to the (lowest) detection limit of 0.05 wt%, have never verified the presence of La in the zirconia grains composing the oxide scale framework. This result was further confirmed by the experiment whose running conditions and outputs are reported, respectively, in Figs. 9 and 10. In fact, tetragonal zirconia was not detected in plate-2 which was subject of a thermal treatment that, in terms of rapid transient, can be more properly compared to arc-jet test T1. The transformation of zirconia phase, which has also larger thermal expansion coefficient (9 ppm/K) and a poorer thermal conductivity (2 W/(m K)) compared to the un-oxidized diboride-based bulk (see Table 2), easily led to cracking and detaching under rapid thermal transient conditions. It follows that the opening of new accesses for inward oxygen diffusion unfortunately poses a serious warning of reliability for potential multi-uses of this material to operate in extreme conditions of temperature.  5.3. Was the presence of LaB6 in the ZrB2-SiC base system beneficial to manage intense aero-thermal heating?  During exposure to arc-jet stream, the ‘‘in situ’’ growing oxide scale covering the hemispheric ZSL10 sample, because of an intrinsic lower thermal conductivity compared to the inner ZrB2-based bulk, transferred only minor percentages of the heat flux delivered  at the surface, and led the exposed surface to reach radiative steady-state temperature approaching 2300 K. Compared to another LaB6-free ZrB2-15%SiC material tested in very similar conditions [34], the present ZSL10 hemisphere reached comparable equilibrium surface temperature when subjected to a less severe thermal load (i.e. 11.4 against 18.5 MJ/kg). A combination of external conditions like equilibrium surface temperature not dropping below 1800 K, and shear stresses along with oxygen partial pressures in the order of 0.2-0.55 kPa and 10-160 Pa, respectively, did not let BSG of any kind to persist over the round sample surface. Borosilicate glasses from a thermo-physical point of view are unstable above 1800 K, and especially in presence of dissociated oxygen and reduced pressure [35], which were running conditions in the arc-jet experiments. The experimental determination here presented of oxidized surfaces lacking of a residual glassy coverage, and also increasingly rougher in the case of samples containing more LaB6 (i.e. ZSL15 and ZSL20), led the authors to conclude that, when the temperatures approach 1800 K under dissociated oxygen and intense shear stresses, the addition of 10-20% LaB6 into a ZrB2-15%SiC matrix did not provide the necessary improvement in oxidation resistance (Fig. 12). As already pointed out, the role of LaB6 was as crucial as unfavorable in breaking up integrity and coherence of the external oxide scale. In this regard, the authors have already documented that, for a very similar formulation lacking of LaB6 and tested in comparable conditions, the persistency of some glassy coating allows the ZrB2- 15%SiC material to manage heat in excess, and thus to limit significantly the adverse consequences of the oxidation [34]. Excessive surface heating was indeed enhanced by the lack of BSG covering the external surface, that led the specimen to behave like a material with both a reduced global emittance (eT) and an higher catalytic surface recombination (c) to dissociated species. Emittance is an important parameter of merit for UHTC since the steady-state temperature they attain via velocity-induced aero-thermal heating is determined, in part, by the efficiency with which they radiate heat away to the colder surroundings. Very recently, total emittance of 0.8 at 1773 K measured on fully dense (pre-oxidized) ZrB2 + 22 vol%SiC + 8 vol%B4C implied that this material efficiently disposed of thermal energy by radiation as long as the glassy oxide coating (persisting over the external surface) dominated the emissive behavior [36]. Such conclusions have been further strengthened thanks to the enlightening work of Marschall [28] who tested the performances of ZrB2-30SiC UHTC in subsonic high enthalpy dissociated air flows: spontaneous temperature jumps seemed associated with the transition in surface chemistry that led to the loss of protective silica glass and substantially increased the chemical component of heat flux delivered to the surface. In the present study, in those regions where zirconia is the only oxidation product, the thermal load (and thus the surface temperature) is enhanced compared to surface areas eventually occupied by BSG (eT = 0.9 [37] and c 6 10\\x002). About c of zirconia, Balat-Pichelin recently documented that the catalytic activity of zirconia may be higher than that of silica glass [38]. Surface catalycity with respect to dissociated air (it was considered the same for Oxygen  Fig. 11. BSE-SEM micrograph of ZSL15 sample after arc-jet testing: a BSG pool (P) at the oxide/un-reacted bulk interface (dashed line) is indicated. The inset points out the position of the magnified view.  Fig. 12. Visual appearance of hemisphere ZSL10, ZSL15 and ZSL20 after arc-jet testing at very similar conditions.  \\x0c', '452  F. Monteverde et al. / Corrosion Science 75 (2013) 443-453  and Nitrogen) was evaluated by comparison between combined CFD and thermal analysis simulations with experimental data. Such comparison indicated that, considering a global emittance eT of 0.6 for a specific total enthalpy of 11.4 MJ/kg, the reacting surface exhibits a catalytic efficiency lower than 10\\x001. A ZrB2-15%SiC formulation free of LaB6 and tested at about 11 MJ/kg in very similar conditions [34] was characterized by a catalycity efficiency of 10\\x002 that is about one order of magnitude higher than that of the ZSL10 composition: according to Marschall [28], the persistency (or not) of residual glassy coating strongly influences the catalytic efficiency of a material in high enthalpy supersonic air flows. The CFD simulations calculated the magnitude of shear stress on the surface which reaches a local maximum of 0.55 kPa in the vicinity of the stagnation point and decreasing to 0.2 kPa downstream the stream length coordinate. The evaluation of this parameter is rather important because may impose more stringent requirements in the final design of a vehicle flying at hypersonic speed. Such shear stress range (0.2-0.55 kPa), in conjunction to temperatures not falling below 1900 K, were sufficient to wipe off the surface by mechanical friction from any BSG which, compared to a pure silica, possesses a lower melting point and an reduced viscosity due to the presence of La, Zr and B in variable amounts. The presence of foreign elements in a silica glass is known to shift its primary network from the stoichiometric high viscous composition. The well-known feature for MB2-SiC ceramics (M = Zr, Hf) of the creation (at high temperatures) of buried zones depleted of SiC [39,40], was not observed. This phenomenon, very often attributed to the active oxidation of SiC, initiates when conditions of oxygen partial pressure reduction from the exterior through a coherent physical external barrier is steadily maintained vs time. In the present case oxygen diffused inward and reached the reacting interfaces without effective delay, i.e. an oxygen activity gradient through the oxide layer was not achieved. In addition, the combined effect of significant percentages of dissociated oxygen diffusing through the oxide scale (constituted by vacancy-rich zirconia grains) enhanced the oxygen diffusivity, and therefore the quantity of oxidants available at the reacting interfaces. Discontinuities of the oxide scale like a gap, a channel and, more generally, a pathway between zirconia grains filled of BSG certainly were another favorite route to transfer oxidants [41,42]. The topic of oxidants moving through silica-rich liquids, although rather complex, is worth of a specific speculation. In addition to molecular diffusion, oxygen transport in a BSG (viscous) medium may also occur by Grotthuss-type oxygen hopping mechanism mediated by network defects, the Si-O-O-B peroxyl linkage or the oxygen deficient center to name two. Li and co-workers [42] proposed that above 1773 K the dominant oxygen carriers in a BSG liquid glass are network defects instead of molecular oxygen O2. Concentrations of dissociated vs molecular oxygen at the model’s surface was calculated to be, respectively, 19% and 1%. Therefore, the mobility advantage of the peroxyl linkage over O2 suggests that arc-jet testing likely leads to faster oxidation than testing in ordinary radiant furnaces at the same temperatures [43]. Also, because the network defects are chemically incorporated into the glass network and thus interact more strongly with solutes than O2, small changes in the BSG chemistry could lead to differences in the oxygen diffusivity controlled by defect trapping/gettering, larger than those simply associated to the better known vehicular diffusion inside a liquid (described by the Stokes-Einstein relation).  6. Conclusions  Fully dense hot-pressed ZrB2 containing 15%vol SiC and increasing amounts of LaB6 (10, 15, and 20 vol%) at expenses of ZrB2 were  oxidized under high enthalpy supersonic air-flows using an arc-jet facility. Transient thermal analysis processed via CFD provided a good agreement between the numerical steady-state temperature and the temperature distribution obtained experimentally through a two-color pyrometer combined with an IR thermo-camera: peak temperatures up to 2300 K (11.4 MJ/kg of specific total enthalpy) were predicted at the stagnation points. During the exposure to arc-jet supersonic air flow the presence of LaB6 into the ZrB2- 15%SiC base system did not end up in the ‘‘in situ’’ growth of an outer solid scale predominantly made of La2Zr2O7 cubic pyrochlore. XRD and SEM-EDS analyses on the oxidized specimens verified the formation of a partly protective oxide scale made of monoclinic zirconia and a (La + Zr)-containing borosilicate glass. Diffusion of oxygen through zirconia was indicated as the slowest mechanism controlling the oxidation rate of the ZrB2-SiC-LaB6 system in such special environment. The presence of LaB6 gave rise to only rare events of phase separation, which did not contribute to an effective slow-down of the diffusion rate of the oxidants through the oxide scale. The contemporary presence of SiC and LaB6 was as crucial as unfavorable in breaking up integrity and coherence of the external oxide scale, and did not bring the necessary improved oxidation resistance to ZrB2 in such extreme conditions.  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Shi, Oxidation behaviour of zirconium diboride-silicon carbide ceramic composites under low oxygen partial pressure, Corros. Sci. 53 (11) (2011) 3742-3746. [41] S.N. Karlsdottir, J.W. Halloran, Formation of oxide scales on zirconium diboride-silicon carbide composites during oxidation: relation of subscale recession to liquid oxide flow, J. Am. Ceram. Soc. 91 (11) (2008) 3652-3658. J. Li, T.J. Lenosky, C.J. Forst, S. Yip, Thermochemical and mechanical stabilities of the oxide scale of ZrB2 + SiC and oxygen transport mechanisms, J. Am. Ceram. Soc. 91 (5) (2008) 1475-1480. [43] A. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D. Ellerby, Oxidation of ZrB2- SiC ceramics under atmospheric and re-entry conditions, Refract. Appl. Trans. 1 (2) (2005) 1-8.  formation  [42]  \\x0c']"
},{
  "_id": 61,
  "PDF": "Effects of oxygen partial pressure and atomic oxygen on the microstructure of oxide scale of ZrB2–SiC composites at 1500°C.pdf",
  "Text": "['Corrosion Science 73 (2013) 44-53  Contents lists available at SciVerse ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effects of oxygen partial pressure and atomic oxygen on the microstructure of oxide scale of ZrB2-SiC composites at 1500 °C  Ning Li a, Ping Hu a, Xinghong Zhang a,⇑ , Yingzhi Liu b, Wenbo Han a  a National Key Laboratory of Science and Technology on Advanced Composites in Special Environments, Harbin Institute of Technology, Harbin 150001, PR China b Beijing Xinghang Mechanical & Electric Equipment Plant, Beijing 100074, PR China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 14 January 2013 Accepted 28 March 2013 Available online 9 April 2013  Keywords:  A. Ceramic matrix composites B. SEM C. Oxidation  1. Introduction  The oxidation behaviour of ZrB2-20 vol.% SiC composites was investigated based on the microstructural evolution of oxide scale under different oxygen partial pressures at 1500 °C, and the similar experiment was performed in atomic oxygen for comparison. The thickness of the oxide scale increases first and then gradually decreases as the pressure decreases, which is strongly dependent on both total pressure and oxygen partial pressure. The atomic oxygen significantly enhances the oxidation of ZrB2-SiC composites, but has little effect on the microstructure of oxide scale. The oxidation mechanism of ZrB2-SiC composites is also discussed in detail.  Ó 2013 Elsevier Ltd. All rights reserved.  Zirconium diboride (ZrB2) based composites, owing to extremely high melting point [1], hardness [2], and retained strength at hightemperatures [3], as well as good chemical inertness [4] and moderate thermal expansion [5], are being considered as the promising candidate materials for aerospace applications. However, the oxidation resistance of ZrB2-based composites is a major problem for the utilization on thermal protection system and other components of hypersonic vehicles, which are subjected to extreme aerodynamic heating during service [6-9]. For this reason, numerous attempts have been made to investigate the oxidation behaviour and mechanism of ZrB2-based composites. The exposure of monolithic ZrB2 in air above \\x18700 °C will produce a scale consisting of ZrO2 and liquid B2O3. It is believed that the relatively high vapour pressure of B2O3 is the essential reason for the failure of monolithic ZrB2 at elevated temperatures [10,11]. The addition of SiC as the second phase opens up new avenues for the improvement of the oxidation resistance. The formation of protective borosilicate glass on the surface can inhibit the diffusion of oxygen into the inner part of the ZrB2-SiC composites. Nevertheless, either elevated temperatures (above \\x181600 °C) or low oxygen partial pressure will significantly accelerate the oxidation process due to active oxidation of SiC [12,13]. The oxidation behaviour of ZrB2-based composites is very complex and is a function of multiple parameters, such as component, temperature, pressure, and so on, all of which influence the processes and products.  ⇑ Corresponding author. Tel./fax: +86 451 86403016. E-mail address: zhangxh@hit.edu.cn (X. Zhang).  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.03.023  Numerous experiments have been reported on the oxidation behaviour of ZrB2-based composites under ambient pressure, and significant progress has been made [14-16]. However, very little attention has been paid to the oxidation of the ZrB2-based composites in atomic oxygen. As well known, during hypersonic flight in the atmosphere, a large amount of oxygen molecules in bow shock are dissociated into atoms, which partially diffuse to the vehicle surface through the boundary layer, and then these oxygen atoms on the surface will recombine into molecules accompanied by the oxidation of wall material. The oxidation behaviour of ZrB2-SiC composites under high enthalpy, low pressure and atomic oxygen conditions is significantly different from that under ambient pressure [17,18]. Gao et al. [19] demonstrated the formation and evolution of zircon during the oxidation of ZrB2-SiC composites at 200 Pa, and pointed out that the main reason for the formation of zircon can be attributed to the active oxidation of SiC. Han et al. [13] also investigated the oxidation behaviour of ZrB2-SiC composites at 1800 °C with oxygen partial pressures of 0.2 and 2 \\x02 10\\x004 atm, respectively. The effects of partial pressure on the microstructure of oxide scale indicated that low oxygen partial pressure was detrimental to the oxidation resistance of the composites because of the active oxidation of SiC. Nevertheless, the quantitative analysis of the oxidation for ZrB2-SiC composites in a wide range of oxygen partial pressure is unavailable in the open literature. In addition, oxygen atom also has a significant influence on thermal protection materials. The degradation of Ni-based alloys and carbon-based materials in dissociated and molecular oxygen has been summarized. The experimental results showed that these materials exhibited high degradation in atomic oxygen [20,21]. Berton et al. [22] investigated the oxidation of SiC under simulated atmospheric  \\x0c', 'N. Li et al. / Corrosion Science 73 (2013) 44-53  45  re-entry conditions. Atomic oxygen significantly affected the oxidation kinetics as well as the crystalline phase of resulting SiO2, and can accelerate the volatilization of silica. However, the transition between active and passive oxidation was not influenced by atomic oxygen. On the contrary, Rosner and Balat et al. [23-25] pointed out the passive oxidation domain was significantly enlarged in atomic oxygen. The understanding of the response of ZrB2-based composites to atomic oxygen lagged behind that of conventional thermal protective materials. Recently, the effect of atomic oxygen adsorption on the oxidation behaviour of ZrB2-ZrC-SiC composites in air was reported, indicating that the atomic oxygen adsorption on the surface enhanced the oxidation of the composites [26]. However, the oxidation mechanism of ZrB2-based composites has not been studied sufficiently, thus the use of these composites in hypersonic vehicles requires a comprehensive understanding of the oxidation behaviour, especially under simulated re-entry conditions. The purpose of this paper is to investigate the microstructural evolution of oxide scale of ZrB2-20 vol.% SiC composites under different oxygen partial pressures (10-60,000 Pa) in molecular and atomic oxygen using a novel microwave-discharge plasma oxidation apparatus. Effects of oxygen partial pressure and atomic oxygen on the oxidation resistance of ZrB2-SiC composites at 1500 °C were determined in terms of the oxide scale thickness as well as oxidation products. The evolution of surface morphology and microstructure of oxide scale were also provided for better understanding of the oxidation mechanism of ZrB2-SiC composites.  2. Materials and methods  2.1. Materials processing  Commercially available ZrB2 (2 lm, >99.5%, Northwest Institute for non-ferrous metal research, China), SiC (0.5 lm, >99.5%, Weifang Kaihua Micro-powder Co., Ltd., China) powders were used as raw materials. The powder mixture of ZrB2 plus 20 vol.% SiC was ball-mixed for 10 h in a polyethylene bottle using ZrO2 balls and ethanol as grinding media. After mixing, the slurry was dried in a rotary evaporator and screened. Then the mixture was hot-pressed at 1950 °C for 1 h under a uniaxial load of 30 MPa in Ar atmosphere. Bulk density and theoretical density were evaluated using the Archimedes method and the rule-of-mixture, respectively. Round-shaped specimens with Ø15 mm \\x02 3 mm were cut from the billet, and then polished to 1 lm finish. Specimens were ultrasonically cleaned in ethanol before testing.  2.2. Oxidation apparatus and tests  The isothermal oxidation of ZrB2-20 vol.% SiC composites was carried out using microwave-discharge plasma apparatus, as illustrated in Fig. 1. The experimental chamber consisted of a quartz tube, 550 mm length and 50 mm inner diameter, with CaF2 viewport at the upper end for pyrometry measurements. The surface temperature of the specimen was measured by a two-colour optical pyrometer (Raytek Ltd., Marathon MR1S, USA) with a measurement range of 1000-3000 °C. The specimen placed on an alumina rod inside the quartz tube was heated to 1500 °C by a high frequency induction device and the surface temperature fluctuation than 5 °C. The fast heating rate (\\x181000 °C/min) of was no larger the high frequency induction device minimized the effects of oxidation during heating process. The mass flow rate of pure oxygen into the experimental chamber was controlled by the flow metre, and the desired pressure values controlled by the pressure regulator and mass flow metre were 60,000, 20,000, 4000, 1000, 100 and 10 Pa, respectively. It should be mentioned that the oxygen partial pressure equaled to the total pressure in the present experiment.  Water out  Water in   Optical pyrometer   CaF2 viewport   Microwave  generator   Mass flow meter   CCD  PC   Wave guide  Quartz tube   O2 gas   Specimen  Optical fiber  Induction coil  Pressure gauge  Alumina rod  Vacuum pump  Fig. 1. Schematic presentation of the plasma oxidation apparatus.  The atomic oxygen was generated by a microwave generator (Nanjing Huiyan Ltd., MY1500S, China) with a frequency of 2.45 GHz and a working power of 1000 W. This novel apparatus can realize the environment of various oxygen partial pressures and atomic oxygen simultaneously.  2.3. Thermodynamic calculation  The experimental study is supplemented by thermodynamic calculations of oxidation reactions under different pressures. Generally, the Gibbs free energy of reaction depends on the Gibbs free energy of formation of reactants and products as given by the relationship  X  X  DG0  r ðT ; p0 Þ ¼  v prodDG0 f ;prod ðT ; p0 Þ \\x00  v reactDG0 f ;react ðT ; p0 Þ  ð1Þ  f ;react ðT ; p0 Þ and DG0  where DG0 r ðT ; p0 Þ is the change of Gibbs free energy for the reaction under standard conditions; DG0 standard Gibbs free energy of formation of reactants and products at the temperature T, respectively. v denotes the stoichiometric coefficient, and p0 is the standard pressure (1 atm).The Gibbs free energy of formation of gas species i, DGf ;i ðT ; pÞ, at the pressure p is calculated by the following equation [27]:  f ;pord ðT ; p0 Þ are the  DGf ;i ðT ; pÞ ¼ DG0 f ;i ðT ; p0 Þ þ RT lnðp=p0 Þ  ð2Þ  where R is the ideal gas constant. In addition, the Gibbs free energy of formation of condensed phase species j, DGf ;j ðT ; pÞ, is assumed to be independent on the pressure, which can be given by:  DGf ;j ðT ; pÞ ¼ DG0  f ;j ðT ; p0 Þ  Therefore, the Gibbs free energy of reaction at can be represented as  X  X  ð3Þ  the pressure p  DGr ðT ; pÞ ¼ ¼ DG0 r ðT ; p0 Þ þ  v prodDGf ;prod ðT ; pÞ \\x00  X  v iRT lnðp=p0 Þ  v reactDGf ;react ðT ; pÞ  ð4Þ  where DGr ðT ; pÞ is the change of Gibbs free energy of reaction at the pressure p; v i is the stoichiometric coefficient of gas species. In the present paper, it is assumed that the gaseous species in reaction is ideal gas. For ZrB2, the oxidation reaction is described as follows,  ZrB2 ðsÞ þ 5=2O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðgÞ  ð5Þ  \\x0c', '46  N. Li et al. / Corrosion Science 73 (2013) 44-53  The Gibbs free energy for this reaction at the pressure p can be represented based on the assumptions and Eq. (4) as  DGr ðT ; pÞ ¼ DG0 r ðT ; p0 Þ þ RT ln  ðpB2 O3 =p0 Þ  ðpO2 =p0 Þ5=2  ð6Þ  where pO2 and pB2 O3 are the pressure of O2 and B2O3, respectively. The Gibbs free energy of reaction is the algebraic sum of the Gibbs free energy of formation of reactants and products. The Gibbs free energy of formation of O2 and B2O3 at the pressure p can be calculated by Eq. (2), which indicates that the O2 pressure and B2O3 pressure in Eq. (6) should be equal to the pressure p, respectively. Therefore, the Gibbs free energy for this reaction at the pressure p can be further rewritten as  DGr ðT ; pÞ ¼ DG0 r ðT ; p0 Þ þ RT lnðp=p0 Þ\\x003=2  ð7Þ  Based on the above-mentioned thermodynamic analysis, the Gibbs free energy of other oxidation reactions for both ZrB2 and SiC with molecular and atomic oxygen can also be calculated using JANAF tables [28].  2.4. Characterization of oxidized specimens  The average thickness of the oxide scale was measured from the fractured surfaces of at least 3 specimens. The morphology of oxidized surface and the microstructure of fractured surface were observed by scanning electron microscopy (SEM, FEI Sirion, Holland) equipped with energy dispersive spectroscopy (EDS, EDAX Inc., USA) for chemical analysis. Crystalline phases of the specimen after oxidation were analyzed by X-ray diffraction (XRD, Rigaku, Japan) at an incidence angel of 2°. The chemical composition of oxide scale was identified using X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific, USA) with monochromatic Al Ka radiation.  3. Results and discussion  The photographs of ZrB2-20 vol.% SiC specimens oxidized at 1500 °C under various oxygen partial pressures for 30 min are presented in Fig. 2. As the oxygen partial pressure decreases from  60,000 to 1000 Pa (Fig. 2a-d), the surface colour changes from dark to white, and then turns to gray under lower oxygen partial pressures, as shown in Fig. 2e and f. The gradual change in colour of the specimen surface can be attributed mainly to the different oxidation products and the surface microstructures. The XRD patterns of specimens oxidized under different oxygen partial pressures are shown in Fig. 3. It is evident that the oxide scale is mainly composed of monoclinic zirconia, and a very weak peak corresponding to stishovite can be identified at 60,000, 20,000 and 4000 Pa, respectively. The XRD patterns are magnified in the inset of Fig. 3 for better comparison, from which a weak and very broad diffraction peak between 15° and 25° should be considered as the amorphous silica-rich glass, and the amorphous peak disappears as the oxygen partial pressure decreases to 1000 Pa and below. Fig. 4 shows the surface SEM images of the oxidized specimens, and the evolution of the surface microstructure can explain the change in macrographs of oxide surfaces. From Fig. 4a, the surface is dark and completely covered with a continuous and smooth silica-rich glass layer after oxidation at 60,000 Pa. As the oxygen partial pressure decreases to 4000 Pa, a large number of white particles mainly composed of ZrO2 are detected beneath the outer glass layer as shown in Fig. 4b and c, indicating that the silica-rich glass layer becomes much thinner. The low total pressure accelerates the evaporation of silica-rich glass, inducing the formation of a thinner silica-rich glass layer [13]. When the oxygen partial pressure is further decreased to 1000 Pa, as shown in Fig. 4d, the surface turns to white, and only zirconia can be identified due to the active oxidation of SiC. The variation of oxidation mechanism between 1000 and 4000 Pa corresponding to passive or active oxidation is consistent with our previous theoretical calculations [11]. Therefore, the difference of surface composition is responsible for the variation of oxide surface colour (from dark to white). In addition, the oxide surfaces are much rougher at 100 and 10 Pa (Fig. 4e and f), which leads to the change in surface from white to gray. The distribution of ZrO2 in surface is consistent with that of ZrB2 in original material. It can be expected that a ZrB2 grain is oxidized in situ to several fine ZrO2 grains, and the microstructural skeleton of the matrix does not  (a)   (b)  (c)  5mm   5mm   5mm   (d)   (e)  (f)  Fig. 2. Photographs of ZrB2-20 vol.% SiC composites oxidized at 1500 °C for 30 min: (a) 60,000 Pa, (b) 20,000 Pa, (c) 4000 Pa, (d) 1000 Pa, (e) 100 Pa and (f) 10 Pa.  5mm   5mm   5mm   \\x0c', 'N. Li et al. / Corrosion Science 73 (2013) 44-53  47  Fig. 5 shows the microstructural evolution of the fractured surfaces of specimens oxidized under different oxygen partial pressures at 1500 °C. As seen from Fig. 5a, a thin and dense oxide scale is observed after oxidation at 60,000 Pa, which consists of two distinct layers: (1) a silica-rich outer layer and (2) a subscale of crystalline zirconia embedded in amorphous silicate. As the pressure decreases from 60,000 to 20,000 Pa, there is no detectable change except for the thickness of oxide scale (Fig. 5b). The average thickness of the oxide scale summarized in Fig. 6 changes from 18 to 24 lm while the silica-rich layer becomes thinner. This can be attributed to the evaporation of silica-rich glass caused by the reduction of total pressure. A comparative test was performed at ambient pressure in air where the oxygen partial pressure is about 20,000 Pa, and the oxide scale of specimen at 1 atm in air (as shown in Fig. 10) is thinner than that at 20,000 Pa. These results show that the total pressure has a great impact on the oxidation resistance of ZrB2-SiC composites. After oxidation at 4000 Pa, the average thickness of the oxide scale increases to 39 lm as shown in Fig. 5c, however, the microstructure of fractured surface of specimen dose not significantly change. Fig. 5d shows SEM micrograph of the fractured surface of the specimen oxidized at 1000 Pa, and provides a wealth of information for the microstructure of oxide scale. The oxide scale becomes porous and loose, and no silica glass is identified in the whole oxide scale. Using SEM analysis  Fig. 3. XRD patterns of ZrB2-20 vol.% SiC composites oxidized under various oxygen partial pressures: (a) 60,000 Pa, (b) 20,000 Pa, (c) 4000 Pa, (d) 1000 Pa, (e) 100 Pa and (f) 10 Pa. The XRD patterns around 20° are magnified in insert showing the change of the amorphous peak.  change. The migration, agglomeration and growth of ZrO2 with liquid products can be neglected at low oxygen partial pressure in the present experiment.  (a)   (c)   (e)   (b)  100µm   (d)  100µm   (f)  100µm   5µm   Fig. 4. Surface morphology of ZrB2-20 vol.% SiC composites oxidized at 1500 °C for 30 min: (a) 60,000 Pa, (b) 20,000 Pa, (c) 4000 Pa, (d) 1000 Pa, (e) 100 Pa and (f) 10 Pa.  5µm   5µm   \\x0c', '48  N. Li et al. / Corrosion Science 73 (2013) 44-53  (a)   (c)  (e)   (b)  10µm   (d)  10µm   porous ZrO2 layer   porous ZrB2 layer   incomplete reaction layer  10µm   10µm   (f)  Fig. 5. Micrographs of fractured surfaces for ZrB2-20 vol.% SiC composites oxidized at 1500 °C for 30 min: (a) 60,000 Pa, (b) 20,000 Pa, (c) 4000 Pa, (d) 1000 Pa, (e) 100 Pa and (f) 10 Pa.  10µm   10µm   combined with EDS, the oxide scale can be identified as three layers: (1) porous ZrO2 layer, (2) porous ZrB2 layer (so-called SiC-depleted layer) and (3) incomplete reaction layer of ZrB2 and SiC. The change in the microstructure of oxide scale exhibits different oxidation mechanisms. The active oxidation of SiC at 1000 Pa is observed, with a rapid increase in the average thickness of the oxide scale to 58 lm. The microstructures of oxide scales at 100 and 10 Pa are similar to that at 1000 Pa, as shown in Fig. 5e and f. Nevertheless, the average thicknesses of the oxide scales of these specimens are decreased to 53 and 42 lm, respectively. These results can be attributed to the low oxygen flux. A significant reduction in the amount of oxygen species available inhibits the growth of oxide scale at very low pressure [29]. Furthermore, some deposited zirconia was observed on the inner side of quartz tube after testing at 10 Pa, indicating that the evaporation of zirconia could not be neglected at very low pressure. Consequently, the microstructure of the oxide scale of specimen at 10 Pa is the most loose in all of specimens. The amorphous silica-rich glass on the surface of the specimen acts as a protective barrier to inhibit the diffusion  Fig. 6. Thicknesses of the oxide scales of ZrB2-20 vol.% SiC composites oxidized under various oxygen partial pressures.  \\x0c', 'N. Li et al. / Corrosion Science 73 (2013) 44-53  49  5mm   5µm   20µm   (a)   (b)   5mm   5µm   20µm   Fig. 7. Photographs and SEM micrographs of ZrB2-20 vol.% SiC composites oxidized in atomic oxygen at 1500 °C for 30 min: (a) 100 Pa; (b) 10 Pa.  of oxygen into the inner part of the matrix [12]. The average thicknesses of silica-rich layer in the oxide scales at 60,000, 20,000 and 4000 Pa are 4.5, 3.3 and 2.6 lm, respectively, and no amorphous silica-rich glass can be identified as the pressure decreases to 1000 Pa and below. It is evident from Fig. 4 that the thickness of silica-rich layer decreases as the total pressure decreases, which is consistent with the analysis of surface morphology. The above results show that the thickness of the oxide scale is affected by both oxygen partial pressure and total pressure. The oxygen partial pressure is responsible for the variation of oxidation reactions and the amount of oxidant, and the total pressure influences the evaporation of oxidation products. In actual flight, thermal protective materials are subjected to extreme thermal and chemical environments containing a large amount of oxygen atoms. It thus appears essential to investigate the oxidation behaviour of ZrB2-SiC composites in dissociated oxygen. As seen from Fig. 7, the specimens oxidized in atomic oxygen are gray, and the surfaces are mainly composed of ZrO2 owing to the active oxidation at low oxygen partial pressure [30]. The surface morphology and the microstructure of the fractured surface for the specimen oxidized in atomic oxygen are not fundamentally different from those in molecular oxygen. In addition, the average thickness of the oxide scale in atomic oxygen changes from 76 to 64 lm as the oxygen partial pressure decreases from 100 to 10 Pa, which is about 1.5 times as that in molecular oxygen. Based on these experimental results, it is suggested that the atomic oxygen significantly enhances the oxidation of ZrB2-SiC composites, but does not change their oxidation behaviour. The oxide scale of specimen oxidized at 10 Pa in atomic oxygen was partially polished to identify chemical composition of the products at the bottom of the scale using XPS. The sub-surface of oxide scale was cleaned by Ar+ ion sputtering for 1 min, and the spectra were calibrated by the C 1s signal detected at 285 eV from adventitious carbon. As shown in Fig. 8a, the peaks at 182.8 eV (Zr 3d5/2) and 185.1 eV (Zr 3d3/2) as the main components correspond to Zr-O bond, and the relative intensity of Zr-B bond in spectra, 178.6 eV (Zr 3d5/2) and 181.0 eV (Zr 3d3/2), are lower than that of Zr-O bond. It is expected that ZrB2 is not completely oxidized into  ZrO2 in the scale. In addition, new peaks at 181.1 eV (Zr 3d5/2) and 183.5 eV (Zr 3d3/2) could be associated with intermediate compounds Zr-O-B during the oxidation. The existence of ZrOxBy has been reported by Hwang et al. [31], however, the systematic studies are rarely reported in the open literature. SiC is identified at 100.6 eV from the corresponding Si 2p spectra (Fig. 8b). Aside from Si-C bond, Si-O and Si-O-C bond are observed at 103.2 and 102.1 eV, respectively. The intermediate components SiOxCy could exist at the interface between SiO2 and SiC [32,33]. It is probable that not all Si-C bonds are replaced by Si-O bonds during oxidation, therefore, the intermediate component is nonstoichiometric. The results of curve fitting C 1s spectra are presented in Fig. 8c. The peak at 285 eV is assigned to C-C bond from adventitious carbon. It should be mentioned that carbon as oxidation product resulting from the active oxidation of SiC would contribute to the C 1s spectra. Other components are mainly related to SiC and SiOxCy located at 283.3 and 286.7 eV, respectively. The large full width at half maximum (FWHM) of the peaks indicates the high chemical inhomogeneity. As shown in the O 1s spectra (Fig. 8d), two main components located at 530.8 and 532.2 eV, respectively, are attributed to ZrO2 and SiO2. The intensity of Si-O-C signal may be very weak and completely masked by the O-Zr and O-Si signals. Therefore, the Si-O-C bond is difficult to be detected in the O 1s spectra compared with that in C 1s and Si 2p spectra. The peak located at 187.5 eV from B 1s spectra (Fig. 8e) is observed and identified as ZrB2. Comparison of specimens after testing in different oxidizing atmospheres suggests that atomic oxygen has high reaction activity for the oxidation of ZrB2-SiC composites. Thermodynamic calculation of reactions for both ZrB2 and SiC in different oxidizing atmospheres as shown in Fig. 9 indicates that these reactions with molecular and atomic oxygen are favoured in the temperature range of 1000-2000 K. It should be noted that the Gibbs free energy of reactions for both ZrB2 and SiC with atomic oxygen are lower than those with molecular oxygen under the same conditions. Consequently, the reactions of both ZrB2 and SiC with atomic oxygen are preferential compared with molecular oxygen. In order to further understand the oxidation mechanism of ZrB2-SiC composites, the oxidation behaviour of SiC should be  \\x0c', '50  N. Li et al. / Corrosion Science 73 (2013) 44-53  Fig. 8. XPS spectra of ZrB2-20 vol.% SiC composites oxidized in atomic oxygen at 1500 °C for 30 min: (a) Zr 3d, (b) Si 2p, (c) C 1s, (d) O 1s and (e) B 1s.  investigated in detail. It has been widely recognized that the oxidation of SiC in ZrB2-based composites is very complex especially at high temperatures, and the oxidation products have a significant influence on the oxidation resistance of these composites [34,35]. The main oxidation reactions of SiC are described as follows,  SiCðsÞ þ 3=2O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ  SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  SiCðsÞ þ 1=2O2 ðgÞ ! SiOðgÞ þ CðsÞ  ð8Þ  ð9Þ  ð10Þ  SiOðgÞ þ 1=2O2 ðgÞ ! SiO2 ðlÞ  CðsÞ þ 1=2O2 ðgÞ ! COðgÞ  ð11Þ  ð12Þ  The oxidation mechanism and resulting products are strongly dependent on the oxidation temperature and oxygen partial pressure [30]. It is worth noting that the oxygen partial pressure and temperature (e.g. for dynamic oxidation) in reaction region are different from those on the surface. According to thermodynamic theory, the oxygen partial pressure beneath the scale should be significantly lower than that on specimen surface due to the oxygen chemical potential gradient associated with any diffusion profile.  \\x0c', 'N. Li et al. / Corrosion Science 73 (2013) 44-53  51  Fig. 9. Variations of Gibbs free energy as a function of temperature for reactions of ZrB2 (a) and SiC (b) in different oxidizing atmospheres.  (a)   (b)  (c)  10µm   0.5µm  Fig. 10. Micrographs and EDS analysis of ZrB2-20 vol.% SiC composites oxidized in air at 1500 °C for 30 min: (a) low-magnification SEM image, (b) high-magnification SEM image and (c) EDS spectrum of partially oxidized SiC grain in (b).  When the ZrB2-SiC composites are exposed to air at high temperatures, SiC is oxidized according to Eq. (8) to produce liquid SiO2 on the surface as a barrier of oxygen diffusion in the initial stage of oxidation process. Subsequently, the thickness of the oxide scale increases as the oxidation time increases. When the thickness of the oxide scale increases to a threshold, the oxygen partial pressure at the interface between the oxide scale and the substrate is low enough to promote the active oxidation of SiC by Eq. (9). The gaseous SiO diffuses outward along the grain boundaries mainly, and reacts with oxygen by Eq. (11) to produce amorphous silica filling the porous skeleton and on the surface of specimen. As the oxygen partial pressure further decreases, carbon is preferentially generated according to Eq. (10), which has been confirmed using TEM [36], and then the gaseous CO is further obtained with increasing oxidation time (as shown in Eq. (12)). Fig. 10a shows the fractured surface of specimen oxidized in air at 1500 °C for 30 min. A thin and dense oxide scale with silica-rich outer layer is identified, and some partially oxidized ZrB2 and SiC grains are found at the bottom of oxide scale (Fig. 10b). EDS analysis (Fig. 10c) shows that the consumption of SiC is attributed mainly to the active oxidation, and only trace oxygen is measured on the surface of shrinking SiC grain. Many partially oxidized SiC grains with various morphologies have been found at the interface between the oxide scale and the substrate, as shown in Fig. 11. The process of active oxidation of SiC phase can be inferred by the morphological evolution of SiC corresponding to the different locations in the scale. From Fig. 11a, the oxidation of SiC grain at the bottom of the oxide scale could be considered as the initial stage. There are many grooves on the surface of SiC grain. It is reasonable to consider that oxygen diffuses to the  grain surface, and then the grain preferentially oxidizes at boundaries, twin boundaries and other defects, which results in the formation of morphology with oriented grooves. With increasing oxidation time, large amounts of ravines can be clearly identified on the surface as shown in Fig. 11b, while the SiC grain still maintains its shape. Fig. 11c shows the morphology of SiC grain in the late stage of oxidation process. SiC grain shrinks significantly, and the surface is filled with lots of fine particles which are mainly composed of SiC confirmed by EDS. The fractured surface of partially oxidized SiC grain is shown in Fig. 11d. There is no other condensed phase except for the SiC on the surface, which is attributed to the active oxidation of SiC. From Fig. 11e, it can be seen that the oxidation rate of SiC is much more than that of ZrB2 beneath the scale, resulting in preferential oxidation of SiC grain. As shown in Fig. 11f, SiC grains are completely depleted and fully removed, leaving the ZrB2 skeleton with large amounts of pores which act as channels for oxygen diffusing into the substrate. It should be noted that large numbers of fine zirconia grains cohere and attach to the surface of channels.  4. Conclusions  In the present work, ZrB2-20 vol.% SiC composites were oxidized under various oxygen partial pressures (10-60,000 Pa) and in different oxidizing atmospheres (molecular oxygen and atomic oxygen) at 1500 °C. The following conclusions can be drawn:  (1) The thickness of the oxide scale is significantly dependent on both total pressure and oxygen partial pressure. Low total pressure accelerates the volatilization of silica-rich glass,  \\x0c', '52  N. Li et al. / Corrosion Science 73 (2013) 44-53  (a)   (c)   (e)   (b)  0.5µm   (d)  0.5µm   (f)  0.5µm   0.5µm   Fig. 11. Micrographs of partially oxidized SiC grains and microstructures in oxide scale of ZrB2-20 vol.% SiC composites after oxidation at 1500 °C.  0.5µm   0.5µm   and low oxygen partial pressure promotes the active oxidation of SiC, which results in a rapid increase in thickness of the oxide scale. The thickness of the oxide scale gradually increases from 18 to 58 lm with decreasing the pressure in the range of 60,000-1000 Pa. Nevertheless, at very low pressure (100 and 10 Pa), the thicknesses of the oxide scales are decreased to 53 and 42 lm, respectively, owing to a significant reduction in the amount of oxygen. (2) Atomic oxygen has high reaction activity and significantly promotes the oxidation of ZrB2-SiC composites. The thickness of the oxide scale in atomic oxygen is about 1.5 times as that in molecular oxygen. However, the oxidation products, surface morphology and the microstructure of oxide scale in atomic oxygen are analogous to those in molecular oxygen. (3) The oxidation of SiC depends strongly on the oxygen partial pressure in reaction region which is different from that on the surface due to the oxygen chemical potential gradient. At the interface between oxide scale and substrate, oxygen  is first transported through grain boundaries, and then reacts with SiC grains in defects preferentially. The SiC grain gradually shrinks due to the active oxidation, leaving the ZrB2 skeleton with large amounts of pores which act as channels for oxygen diffusing into the substrate.  Acknowledgements  This work was financially supported by the National Science Foundation of China (51072042 and 50972029) and the Fundamental Research Funds for the Central Universities (HIT.NSRIF. 2012030).  References  [1] F. Monteverde, The thermal stability in air of hot-pressed diboride matrix composites for uses at ultra-high temperatures, Corros. Sci. 47 (2005) 2020- 2033.  \\x0c', 'N. Li et al. / Corrosion Science 73 (2013) 44-53  53  [2] T. Zhu, W.J. 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Meng, Oxidation behavior of zirconium diboride-silicon carbide at 1800 °C, Scripta Mater. 57 (2007) 825-828. [14] T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, A model for the oxidation of ZrB2, HfB2 and TiB2, Acta Mater. 55 (17) (2007) 5999-6010. [15] S.B. Zhou, W.J. Li, P. Hu, C.Q. Hong, L. Weng, Ablation behavior of ZrB2-SiC- ZrO2 ceramic composites by means of the oxyacetylene torch, Corros. Sci. 51 (2009) 2071-2079. [16] Z.J. Wu, Z. Wang, Q. Qu, G.D. Shi, Oxidation mechanism of a ZrB2-SiC-ZrC ceramic heated through high frequency induction at 1600 °C, Corros. Sci. 53 (2011) 2344-2349. J.L. Vérant, N. Perron, O. Gerasimova, M. Balat-Pichelin, V. Sakharov, A. Kolesnikov, O. Chazot, P. Omaly, Microscopic and macroscopic analysis for TPS SiC material under earth and mars reentry conditions, AIAA 2006-7947. [18] M. Balat, M. Czerniak, J.M. Badie, Thermal and chemical approaches for oxygen catalytic recombination evaluation on ceramic matericals at high temperature, Appl. Surf. Sci. 120 (1997) 225-238. [19] D. Gao, Y. Zhang, J.Y. Fu, C.L. Xu, Y. Song, X.B. Shi, Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: formation of zircon, Corros. Sci. 52 (2010) 3297-3303.  J. Am. Ceram. Soc. 88 (12)  formation  [13]  [17]  in dissociated  space shuttle TPS metals  [20] W.P. Gilbreath, The degradation of oxygen, AIAA 1972-262. [21] K. Fujimoto, T. Shioya, K. Satoh, Degradation of carbon-based materials due to impact of high-energy atomic oxygen, Int. J. Impact Eng. 28 (2003) 1-11. [22] B. Berton, M.P. Bacos, D. Demange, J. Lahaye, High-temperature oxidation of silicon carbide in simulated atmospheric re-entry conditions, J. Mater. Sci. 27 (1992) 3206-3210. [23] D.E. Rosner, H.D. Allendorf, High temperature kinetics of the oxidation and nitridation of pyrolytic silicon carbide in dissociated gases, J. Phys. Chem. 74 (9) (1970) 1829-1839. [24] M. Balat, G. Flamant, G. Male, G. Pichelin, Active to passive transition in the oxidation of silicon carbide at high temperature and low pressure in molecular and atomic oxygen, J. Mater. Sci. 27 (1992) 697-703. [25] M.J.H. Balat, Determination of the active-to-passive transition in the oxidation of silicon carbide in standard and microwave-excited air, J. Eur. Ceram. Soc. 16 (1996) 55-62. [26] D. Gao, Y. Zhanga, C.L. Xu, Y. Song, X.B. Shi, Atomic oxygen adsorption and its effect on the oxidation behaviour of ZrB2-ZrC-SiC in air, Mater. Chem. Phys. 126 (2011) 156-161. [27] D.A. Holmes, B.K. Harrlson, Estimation of Gibbs free energies of formation for polychlorinated biphenyls, Environ. Sci. Technol. 27 (1993) 725-731. [28] M.W. Chase Jr., NIST-JANAF Thermochemical Table, fourth ed., American Institute of Physics, Woodbury, New York, 1998. J. Marschall, D.A. Pejakovic´ , W.G. Fahrenholtz, G.E. Hilmas, S. Zhu, J. Ridge, D.G. Fletcher, C.O. Asma, J. Thömel, Oxidation of ZrB2-SiC ultrahigh-temperature ceramic composites in dissociated air, J. Thermophys. Heat Transfer 23 (2) (2009) 267-278. [30] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Oxidation of zirconium diboridesilicon carbide at 1500 °C at a low partial pressure of oxygen, J. Am. Ceram. Soc. 89 (10) (2006) 3240-3245. [31] S.S. Hwang, A.L. Vasiliev, N.P. Padture, Improved processing and oxidationresistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids, Mater. Sci. Eng. A 464 (2007) 216-224. [32] L. Scatteia, R. Borrelli, G. Cosentino, E. Bêche, J.L. Sans, M. Balat-Pichelin, Catalytic and radiative behaviors of ZrB2-SiC ultrahigh temperature ceramic composites, J. Spacecraft Rockets 43 (5) (2006) 1004-1012. [33] A. Avila, I. Montero, L. Galán, J.M. Ripalda, R. Levy, Behavior of oxygen doped SiC thin films: an x-ray photoelectron spectroscopy study, J. Appl. Phys. 89 (1) (2001) 212-216. [34] X.H. Zhang, P. Hu, J.C. Han, Structure evolution of oxidation in air, J. Mater. Res. 23 (7) (2008) 1961-1972. [35] W.B. Han, P. Hu, X.H. Zhang, J.C. Han, S.H. Meng, High-temperature oxidation at 1900 °C of ZrB2-xSiC ultrahigh-temperature ceramic composites, J. Am. Ceram. Soc. 91 (10) (2008) 3328-3334. [36] D.W. Ni, G.J. Zhang, F.F. Xu, W.M. Guo, Initial stage of oxidation process and microstructure analysis of HfB2-20 vol.% SiC composite at 1500 °C, Scripta Mater. 64 (2011) 617-620.  ZrB2-SiC  [29]  during  the  \\x0c']"
},{
  "_id": 62,
  "PDF": "Effects of SiC on Oxidation Behavior of ZrB2-Based Composites with MoSi2 and SiC Additives.pdf",
  "Text": "['Key Engineering Materials Vols. 602-603 (2014) pp 457-462 Online available since 2014/Mar/12 at www.scientific.net © (2014) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.602-603.457  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (ID: 129.93.16.3, University of Nebraska-Lincoln, Lincoln, USA-02/02/15,15:53:45)  Effects of SiC on Oxidation Behavior of ZrB2-Based Composites with MoSi2 and SiC Additives Shuqi Guo Hybrid Materials Unit, National Institute for Materials Science,  1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan GUO.Shuqi@nims.go.jp Keywords: Zirconium diboride; Molybdenum disilicide; Silicon carbide;  Oxidation behavior Abstract. In this study, oxidation behavior of ZrB2-MoSi2-SiC composite was investigated in the hot-pressed 5-20 vol% SiC-containing ZrB2-20 vol% MoSi2-based composites which were exposed to dry air between 1100°C and 1500°C up to 10 hours. The effects of SiC additive on the oxidation behavior were assessed. Experimental results showed that the weight gain due to oxidation exposure in air increased with increasing exposure temperature and exposure time. Parabolic oxidation behavior was observed for all the compositions composites. On the other hand, the weight gain decreased with increasing amount of SiC added. The addition of SiC improved the oxidation resistance of the composites, and the improvement was enhanced with increasing amount of SiC added. In addition, X-ray diffraction was used to identify major crystalline phases present in both the as-received and the post-oxidized composites. The oxidized sample surface was characterized by scanning electron microscopy and energy-dispersive X-ray spectroscopy. The microstructure of the post-oxidized composites consisted of two characteristic regions: oxidized reactive region and unreactive bulk material region. Furthermore, the oxidized reactive region divided into an outermost dense silica-rich scale layer and oxidized reactive mixture layer. The improvement of the oxidation resistance due to the addition of SiC is associated with the presence of the thicker dense outermost scale layer which inhibited inward diffusion of oxygen through it. Introduction Zirconium diboride (ZrB2)-based ceramic composites have become an important class of potential candidate materials for various ultra-high temperature structural applications because they have an extremely high melting point (>3000°C), high thermal and electrical conductivities, chemical inertness against molten metals and good thermal shock resistance [1,2]. A major problem with ZrB2-based ceramics is the poor oxidation resistance at high temperature where they are considered to be applied as structural materials in ultra-high temperature oxidizing environments [3,4]. Heating ZrB2 in air produces a scale composed of ZrO2 and B2O3, but the B2O3 has a high vapor pressure and is vaporized above 1300°C [5-9], led to low oxidation resistance. Thus, the oxidation resistance of ZrB2 ceramics must be improved for structural applications in oxidizing environments at a temperature equal to or higher than 1300°C. Si-containing additives such as SiC and MoSi2 are added to ZrB2 [5-12] in order to improve oxidation resistance. These additives led to form a protective borosilicate glass layer at elevated temperature that enhances the oxidation resistance of ZrB2. In particular, ZrB2-MoSi2-SiC system showed the better oxidation resistance compared to ZrB2-MoSi2 or ZrB2-SiC systems, as a result of the complex effects of MoSi2 and SiC [11,12]. However, it is not well-known that effects of temperature and compositions on the oxidation behavior of the ZrB2-MoSi2-SiC system. Generally, the oxidation resistance depended on compositions, exposure temperature and exposure time at a particular temperature, as well as on oxygen content in oxidizing atmosphere. Thus, it could be expected that the oxidation resistance of the ZrB2-MoSi2-SiC system is closely linked with exposure temperature and compositions of materials. In this study, the hot-pressed compacts of 5, 10, 20 vol% SiC-modified 20 vol% MoSi2-ZrB2 composites were exposed to dry air at different temperatures \\x0c', '458  High-Performance Ceramics VIII   between 1100°C and 1500°C for up to 10 h. The oxidation behavior of the composites was examined. X-ray diffraction was used to identify the oxidation products. Morphologies of the post-oxidized samples were observed by scanning electron microscopy. Also, the effects of exposure temperature and amount of SiC added on the oxidation behavior were discussed. Experimental The material used in this study was 20 vol% MoSi2-containing ZrB2-based ceramic composites with SiC additives. In order to examine the oxidation resistance of the composites as well as to learn the effect of amount of SiC added, three compositions of 5, 10 and 20 vol% SiC-modified 20 vol% MoSi2-ZrB2 composites were hot-pressed at 1800°C and 30 MPa for 30 min in vacuum. Fig. 1 shows an example of the backscattered electron FE-SEM micrographs of the ZrB2-MoSi2-SiC composites. The light-grey phase and the intermediate-grey phase in the backscattered electron FE-SEM images were ZrB2 phase and MoSi2 phase, respectively, and the phase with the darkest contrast was SiC phase. The general microstructure of the composites was similar for the three compositions and consisted of the equiaxed ZrB2, MoSi2 and SiC grains. In addition, the SiC particles were randomly dispersed at the ZrB2 and MoSi2 grain junctions. The detailed sintering process and microstructure observations had been reported elsewhere [13]. Hereafter, the three compositions materials are denoted as ZMS5, ZMS10 and ZMS20, respectively. Test specimens with average dimensions of 5 mm × 2.5 mm × 2 mm were cut from the composites plates with a diamond wafering blade. After the specimens were polished with a diamond paste down to 1.0 µm, they were ultrasonically cleaned in acetone and then were kept in an oven at a constant temperature of 100°C prior to oxidation. Oxidation exposure was performed using an electronic furnace in dry air at different temperatures between 1100°C and 1500°C for up to 10 h. Before and after the oxidation, the specimens were weighed, respectively, using an analytical balance with an accuracy of 0.1 mg. X-ray diffraction (XRD) was used to identify the crystalline phases present in the post-oxidized samples. The oxidized surfaces were characterized by scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDX). In addition, in order to examine the oxidation evolution, the oxidized specimens were cut in half, and one of the parts was mounted in epoxy and then was carefully polished with a diamond paste down to 1.0 µm. After the polishing, the SEM observations and EDX analysis were performed on the polished surfaces. Weight Gain, W (mg/cm2)Time at Temperature, t (h) 12ZMS5ZMS10ZMS201500°C1500°C1500°C1100°C1300°C1300°C0246810120246810 2θ(degree)Intensity (a.u.)20304050607080ZrB2ZrO2ZMS10ZMS5ZMS20MoBMoSi2ZrSiO4 Fig. 2 Plots of weight gain as a function of exposure time for the composites. Fig. 3 XRD patterns of the post-oxidized composites at 1500°C for 10 h. \\x0c', 'Key Engineering Materials Vols. 602-603  459   Results and discussion Weight gain. Plots of weight gain as a function of exposure time for the three composition composites are shown in Fig. 2. All three compositions materials showed similar oxidation behaviors: the specific weight increases rapidly within initial 1 h of exposure and then the weight increased gradually with increasing exposure time, independent of the SiC amount and exposure temperature, showing the same oxidation mechanism. In addition, the weight gain decreased with increasing amount of SiC added. Furthermore, the weight gains are lower than those obtained in the SiC-free ZrB2-based composites with MoSi2 exposed at the same oxidation conditions which were previously reported by authors elsewhere [11]. This indicated that oxidation resistance of the composites was improved due to addition of SiC, and this improvement was enhanced with increasing SiC content. At 1300°C or above, in particular, the oxidation resistance was significantly improved. Microstructure of post-oxidized samples. XRD patterns of the post-oxidized samples were showed in Fig. 3. After the oxidation, the primary oxidized phase of ZrO2 is detected in all samples. The minor oxidized phases of MoB and ZrSiO4 phases were present as well. In addition, a trace of amount of unreactived ZrB2 and MoSi2 phases were present. The peak intensity of ZrB2 and MoSi2 phases increased with increasing amount of SiC added, showing improved oxidation resistance, because those signals are from the bulk material beneath the scale layer formed during exposure. This agrees with the observed decrease in weight gain with increasing amount of SiC (Fig. 2). However, the peak of SiC phase is absent in the post-oxidized samples. This indicated that SiC is thoroughly oxidized during the exposure and the residual SiC concentration could not detected. It is known that ZrB2, MoSi2 and SiC phases are oxidized to form ZrO2, B2O3, SiO2 and MoB when they are exposed to dry air at high temperature, as a result of the outward diffusion of constituent elements cations from the bulk to the oxidized surface and the inward diffusion of O through the scale layer. Based on the previous studies [8-12], when the ZrB2, MoSi2 and SiC were exposed to air at high temperature, the following reactions should be occurred. ZrB2(s) + 2.5O2(g) = ZrO2(s) + B2O3(g)  (1) SiC(s) + 1.5O2(g) = SiO2(s) + CO(g) (2)  \\x0c', '460  High-Performance Ceramics VIII   MoSi2(s) + 3.5O2(g) = MoO3(g) + 2SiO2(s) (3) ZrB2(s) + 2MoSi2(s) + 5O2(g) = ZrO2(s) + 2MoB(s) + 4SiO2(s)  (4) In the present study, SiO2, ZrO2, MoB and ZrSiO4 were detected in the post-oxidized samples (Fig. 3). Thus, it could be expected that similar reactions occurred during the exposure at high temperature for the materials investigated in this study. The presence of ZrSiO4 is a result of the reaction between ZrO2 and amorphous silica [12]. Fig. 4 shows the typical examples of SEM images of the surface morphologies and the corresponding EDX analysis results for the oxidized samples. After the oxidation, the surfaces are covered with a continuous silica layer where white ZrO2 particles are embedded, independent on the amount of SiC added. For the post-oxidized sample at 1100°C, the ZrO2 particles did not coarsen significantly. For comparison, after the oxidation at 1500°C the particles significantly coarsened to  1500°C/10hMoBSiZrOZMS20III\\x0c', 'Key Engineering Materials Vols. 602-603  461   form some larger nodules. No cracking behavior was observed in the post-oxidized samples, however. In addition, EDX analysis revealed that the background consisted of only Si and O. This indicated that the background is an amorphous silicate. In contrast, Si, O and Zr were detected in the nodules. This suggests that the nodules particles were an oxide cluster where some larger ZrO2 embedded in a SiO2-rich glass matrix. Fig. 5 shows SEM image of the cross-section and corresponding EDX mapping for the post-oxidized samples at 1100°C for 10 h. The cross-section of the post-oxidized samples is divided into the oxidized reactive region and the unreactive bulk material region. The thickness of the reactive region increased with increasing exposure time, but it became thinner with increasing amount of SiC added. In addition, the reactive region consisted of two different characteristic layers, I and II. Layer I is a dense scale layer and it is rich Si and O. Layer II is SiC-depletion, suggesting that SiC in layer II was oxidized due to outward diffusion of Si and inward diffusion of O. This diffusion led to a formation of the outermost scale on the post-oxidized samples. In addition, increasing exposure temperature resulted in a thicker scale layer and an extensive SiC-depleted layer, as shown in Fig. 6. It is evident that the outermost scale layer was denser and the SiC-depleted layer was more serious after oxidation at 1500°C than after oxidation at 1100°C. Furthermore, only the trace amount of SiC was detected in layer II of the sample oxidized at 1500°C compared to the sample oxidized at 1100°C (Fig. 5). This suggests that the inward diffusion of oxygen and outward diffusion of silicon occurred during oxidation exposure were accelerated with increasing exposure temperature. This means that oxidation mechanism of ZrB2-based composites is mostly dominated by diffusion of Si and O. Summary Remarks The oxidation behavior of the ZrB2-MoSi2-SiC composites was investigated through thermal exposure to dry air between 1100°C and 1500°C up to 10 h. The effects of amount of SiC added on the oxidation behavior of the composites were examined. The major results obtained are drawn as follows: (i) parabolic oxidation behavior was observed for the composites; (ii) the oxidation behavior of the composites was dominated by the outward diffusion of constituent elements cations from the bulk to the oxidized surface and the inward diffusion of O through the scale layer; (iii) the oxidation resistance of the composites was improved due to addition of SiC and the improvement was enhanced with increasing amount SiC added; (iv) the microstructure of the post-oxidized composites consisted of an outermost dense glassy scale layer, middle SiC-depletion layer, and unreactive bulk material, independent on exposure temperature and amount of SiC added; and (v) main oxidation products consisted of ZrO2, MoB, SiO2 and ZrSiO4. Acknowledgments Author would like to thank Dr. T. Mizuguchi, National Institute for Materials Science, for his assistance with SEM observations and EDX analysis. References [1] K. Upadhya, J.-M. Yang and W.P. Hoffmann, Materials for ultrahigh temperature structural applications, Am. Ceram. Soc. Bull. 76 (1997) 51-56. [2] S.Q. Guo, Densification of ZrB2-based composites and their mechanical and physical properties: A review, J. Euro. Ceram. Soc. 29 (2009) 995-1011. [3] E. Wuchina, E. Opila, M. Opeka, W. Fahrenholtz and I. Talmy, UHTCs: Ultra-high temperature ceramic materials for extreme environment applications, Interface, 16 (2007) 30-36. [4] A. Paul, D.D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. Binner, B. Vaidhyanathan, A. Heaton, P. Brown and W.E. Lee, UHTS composites for hypersonic applications, Am. Ceram. Soc. Bull. 91 (2012) 22-29. \\x0c', '462  High-Performance Ceramics VIII   [5] A.K. Kuriakose and J.L. Margrave, Oxidation kinetics of zirconium diboride and zirconium carbides at high temperatures, J. Electrochem. Soc. 111 (1964) 827-831. [6] J.B. Berkowitz-Mattuck, High temperature oxidation, J. Electrochem. Soc. 113 (1966) 908-914. [7] W.C. Tripp, H.H. Davis and H.C. Graham, Effect of an SiC addition on the oxidation of ZrB2, Am. Ceram. Soc. Bull. 52 (1973) 612-616. [8] F. Monteverde and A. Bellosi, Oxidation of ZrB2-based ceramics in dry air, J. Electrochem. Soc. 150 (2003) B552-B559. [9] A. Rezaie, W.G. Fahrenholtz and G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500°C, J. Euro. Ceram. Soc. 27 (2007) 2495-2501. [10] D. Sciti, M. Brach and A. Bellosi, Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite, J. Mater. Res. 20 (2005) 922-930. [11] S.Q. Guo, T. Mizuguchi, T. Aoyagi, T. Kimura and Y. Kagawa, Quantitative electron microprobe characterizations of oxidized ZrB2 containing MoSi2 additives, Oxid. Met. 72 (2009) 335-345. [12] S.Q. Guo, T. Mizuguchi, M. Ikegami and Y. Kagawa, Oxidation behavior of ZrB2-MoSi2-SiC composites in air at 1500°C, Ceram. Int. 37(2011) 585-591.  [13] S.Q. Guo, T. Nishimura, H. Tanaka and Y. Kagawa, Thermal and electrical properties in hot-pressed ZrB2-MoSi2-SiC composites, J. Am. Ceram. Soc. 90 (2007) 2255-2258. \\x0c', 'High-Performance Ceramics VIII   10.4028/www.scientific.net/KEM.602-603   Effects of SiC on Oxidation Behavior of ZrB2-Based Composites with MoSi2 and SiC Additives   10.4028/www.scientific.net/KEM.602-603.457   DOI References  [2] S.Q. Guo, Densification of ZrB2-based composites and their mechanical and physical properties: A  review, J. Euro. Ceram. Soc. 29 (2009) 995-1011.  http://dx.doi.org/10.1016/j.jeurceramsoc.2008.11.008  [9] A. Rezaie, W.G. Fahrenholtz and G.E. Hilmas, Evolution of structure during the oxidation of zirconium  diboride-silicon carbide in air up to 1500°C, J. Euro. Ceram. Soc. 27 (2007) 2495-2501.  http://dx.doi.org/10.1016/j.jeurceramsoc.2006.10.012  [10] D. Sciti, M. Brach and A. Bellosi, Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic  composite, J. Mater. Res. 20 (2005) 922-930.  http://dx.doi.org/10.1557/JMR.2005.0111  [11] S.Q. Guo, T. Mizuguchi, T. Aoyagi, T. Kimura and Y. Kagawa, Quantitative electron microprobe  characterizations of oxidized ZrB2 containing MoSi2 additives, Oxid. Met. 72 (2009) 335-345.        \\x0c']"
},{
  "_id": 63,
  "PDF": "Effects of temperature and the incorporation of W on the oxidation of ZrB2 ceramics.pdf",
  "Text": "['Corrosion Science 80 (2014) 221-228  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effects of temperature and the incorporation of W on the oxidation of ZrB2 ceramics  ⇑  M. Kazemzadeh Dehdashti  , W.G. Fahrenholtz, G.E. Hilmas  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65401, United States  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 22 July 2013 Accepted 24 November 2013 Available online 1 December 2013  Keywords:  A. Ceramic B. SEM, weight loss C. Oxidation  1. Introduction  The effects of tungsten additions and temperature on the oxidation behavior of nominally pure ZrB2 and ZrB2 containing 4, 6 or 8 mol% of W after oxidation at temperatures ranging from 800 to 1600 °C were investigated. For pure ZrB2, the protective liquid/glassy layer covering the surface as a result of oxidation was evaporated above 1500 °C. For (Zr,W)B2 specimens, the liquid/glassy layer was present after exposure up to 1600 °C. The higher stability of the liquid/glassy phase in the W-containing compositions was attributed to the presence of tungsten in the liquid/glassy phase, resulting in improved oxidation resistance for ZrB2 samples containing W.  Ó 2013 Elsevier Ltd. All rights reserved.  Zirconium diboride has advantages over other candidates for hypersonic aerospace applications due to its high melting point (\\x183250 °C) and low theoretical density (6.09 g/cm3) combined with a thermal conductivity as high as 100 W/m K at room temperature [1,2]. Ultrahigh-temperature stability and high thermal conductivity provide ZrB2 with the ability to transfer heat away from the hottest areas of structures and redistribute it to cooler areas, which makes it attractive for sharp leading edges for hypersonic aerospace vehicles [3]. The use of ZrB2-based ceramics for applications requiring elevated temperatures in air is restricted by its oxidation behavior. Assuming stoichiometric oxidation, exposure of ZrB2 to air results in the formation of B2O3 and ZrO2, which leads to measurable mass gain (Eq. (1)) [4].  ZrB2ðsÞ þ 5=2 O2ðgÞ ! ZrO2ðsÞ þ B2O3ðlÞ  ð1Þ  \\x0e  DG  The oxidation reaction is favorable at all temperatures with rxn ¼ \\x001977 þ 0:361T ðkJÞ [5]. Previous studies have divided the oxidation behavior of ZrB2 into three different temperature regimes, although the transition temperatures depend on parameters such as heating rate, gas flow rate, and oxygen activity. The low temperature regime occurs below about 1100 °C. At these temperatures, ZrO2 and B2O3 form a continuous oxide layer that provides passive oxidation protection [6-9]. Analysis concluded that the transport of oxygen through B2O3 was the rate limiting step for oxidation in this regime, which results in parabolic (diffusion-limited growth)  ⇑ Corresponding author. Tel.: +1 573 578 2853; fax: +1 573 341 6934. E-mail address: mk7y8@mst.edu (M. Kazemzadeh Dehdashti).  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.11.030  kinetics for mass gain and changes in the oxide layer thickness [9,10]. Between \\x181100 °C and \\x181400 °C, paralinear kinetics have been observed. In this intermediate temperature regime, the weight change reflects a combination of mass gain from formation of ZrO2 and B2O3, and mass loss due to the evaporation of B2O3 [6,10,11]. Specimens continue to gain mass as the mass of ZrO2 plus retained B2O3 is greater than the mass of diboride reacted plus the mass of volatilized B2O3. As B2O3 evaporates, a porous ZrO2 layer is left behind. In the highest temperature regime, \\x181500 °C and above, the B2O3 is no longer protective and mass gain kinetics become linear due to the non-protective nature of the ZrO2 layer [9]. Parthasarathy et al. [12,13] presented a physical model to predict the oxidation behavior of refractory diborides in these three different temperature regimes. The model assumed that diffusion of dissolved oxygen through boria, in capillaries between nearly columnar blocks of the ZrO2 was the rate-limiting step when condensed boria was present. After evaporation of boria, the oxidation rate was limited by Knudsen diffusion of molecular oxygen through the capillaries. The model agreed with experimental results in predicting weight change, recession, oxide scale thickness and the temperature dependence of the parabolic rate constant for ZrB2. Addition of Si-containing compounds such as SiC [3,4,14-23], MoSi2 [24-26], or TaSi2 [27,28] is the conventional approach to improving the oxidation resistance of diboride ceramics. Improved oxidation resistance is obtained as a result of formation of a borosilicate glassy layer on the surface of the diboride with higher stability compared to the borate glassy layer which forms on nominally pure ZrB2 [4,23]. However, SiC or other Si-containing additives may not be the best choice due to problems, such as rupture of the protective glassy layer as a result of SiO(g) formation [29], or formation of a SiC-depleted layer. Either of these issues has been shown to result in loss of protection at temperatures above 1600 °C [30,31].  \\x0c', '222  M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  Transition metal additives can also improve the oxidation behavior of diborides. Led by Talmy et al., several groups have investigated transition metal additions including Cr-, Ti-, Nb-, V-, and Ta-borides, to improve the oxidation resistance of ZrB2-based composites [31-35]. In addition, Zhang et al. [11,36] reported that WC additions improved the oxidation resistance of ZrB2 ceramics due to the formation of WO3 in the oxide scale, which resulted in liquid phase sintering of ZrO2 and consequently increased the relative density of the scale. The purpose of the present paper was to study the effect of the amount of W added on the oxidation behavior of ZrB2 at temperatures from 800 to 1600 °C. The initial results indicated the effectiveness of W on increasing the stability of the protective liquid/glassy B2O3 scale. Hence, B2O3 glasses with 0, 4, 6, or 8 mol% WO3 were also prepared to investigate the effect of WO3 on weight loss of B2O3 glasses. This study was intended as an initial investigation of oxidation behavior in a controlled furnace oxidation environment with the goal of gaining insight into fundamental oxidation mechanisms prior to testing in more complex environments that are representative of hypersonic flight conditions.  2. Experimental procedure  Specimens for this study were prepared from high purity (>99%) ZrB2 powder (\\x182 lm, Grade B, H.C. Starck, Newton, MA). 2 wt% B4C (\\x180.8 lm, Grade HS, H.C. Starck) was added to all batches to remove the oxide impurities from the powder particle surfaces and enhance densification. For some batches, 4, 6, or 8 mol% tungsten (\\x183 lm, Alfa Aesar, Ward Hill, was added in form of W powder MA) with a reported purity of 99.9%. Hereafter, ZrB2 ceramics containing only B4C are referred to as ‘‘nominally pure ZrB2’’ while ZrB2 with B4C and W additions are referred to as (Zr,W)B2. The powders were mixed by ball milling in butan-2-one (methyl ethyl ketone) with ZrO2 media for 24 h. Measurement of the mass of media before and after milling indicated that the amount of zirconia contamination added to the batches was less than 1 wt% based on the mass of ZrB2 powder. After drying and sieving (\\x0080 mesh), the powders were densified by hot pressing (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 2100 °C for 45 min at a pressure of 32 MPa. For oxidation studies and initial characterization of the as processed microstructures, bars with dimensions of 10 mm by 4 mm by 4 mm were diced from the produced billets. Prior to testing, bars were finished by polishing on all sides using successively finer diamond abrasives with a final polishing step using a 15 lm diamond slurry. The microstructures of the polished surfaces of as processed (Zr,W)B2 and ZrB2 were studied using images obtained by scanning electron microscopy (SEM; S-4700, Hitachi, Japan). Image analysis software (ImageJ, US National Institutes of Health, Bethesda, MD) was used to calculate the amount of B4C remaining after densification. The Archimedes technique, with water as the immersing medium, was used to measure the bulk density of the hot pressed billets. Oxidation studies were performed in a MoSi2 resistance-heated horizontal tube furnace (Model 0000543 Rapid Temperature Furnace, CM Inc., Bloomfield, NJ) equipped with a high-purity alumina  tube with a diameter of 6.35 cm. Specimens were placed on a ridged zirconia setter that was on an alumina D-tube in the center of an alumina tube that was sealed with gas-tight end caps. Specimens were heated at \\x185 °C/min to temperatures ranging from 800 to 1600 °C in flowing air with a flow rate of 0.2 cm/s (linear flow rate was calculated from the volumetric flow rate and the diameter of the tube). Specimens were air quenched to room temperature by removing them from the furnace after the desired oxidation time to minimize changes such as further oxidation that may occur during cooling. The weight and surface area of the specimens were measured before and after oxidation to calculate the weight gain as a result of oxidation. The fracture surfaces observed in the SEM were used to study the microstructure and calculate the thicknesses of the resulting oxide layers. In addition, chemical compositions of the scales were analyzed using energy dispersive spectroscopy (EDS; EDAX, Mahwah, NJ). Borate glasses were prepared from high purity (P99.5%) boric acid powder (H3BO3, ACS reagent, Sigma-Aldrich, Saint Louis, MO) and WO3 (99.99%, metal basis, Alfa Aesar, Ward Hill, MA). The powders were mixed using dry ball milling followed by grinding with a mortar and pestle. The mixtures were calcined at 500 °C in a platinum crucible until no bubbling was observed. The glasses were then melted at 900 °C for 1 h followed by quenching onto copper plates. After grinding and sieving (\\x0080 mesh), thermogravimetric analysis (Netzsch Simultaneous TGA/DTA, Selb, Germany) was performed at temperatures up to 1500 °C using a heating rate of 10 °C/min and a flowing nitrogen atmosphere. Before initiating the analysis, the samples were held for 1 h at 170 °C to eliminate any hydration of the B2O3 that may have occurred during preparation or storage.  3. Results and discussion  All of the compositions reached nearly full density during hot pressing. The measured bulk densities of the specimens are shown in Table 1. To determine the theoretical densities, the area fraction of B4C remaining in each composition after densification was estimated by analyzing SEM images and assumed to be equivalent to the B4C volume fraction. Using true densities of 6.09 g/cm3 for ZrB2, 19.25 g/cm3 for W, and 2.52 g/cm3 for B4C, a volumetric rule of mixtures was used to calculate the theoretical density of nominally pure ZrB2 and (Zr,W)B2 specimens. Combined with the bulk density measurements, all of the specimens had relative densities that were >98%. In addition, the Archimedes’ measurements indicated that open porosity was not significant in any of the compositions. Therefore, porosity should not have an effect on the oxidation behavior. Oxidation of (Zr,W)B2 should produce ZrO2, B2O3, and WO3, assuming that oxidation proceeds stoichiometrically (Eq. (2)), although oxidation is minimal below \\x18800 °C. Individually, the stable phases of these compounds in this temperature range are solid ZrO2 (monoclinic below 1170 °C and tetragonal above), liquid B2O3 (melting temperature \\x18450 °C), and WO3. When combined, the scale formed on the surface of (Zr,W)B2 specimens after oxidation at elevated temperature and cooling to room temperature  Table 1 Area fraction of B4C remaining after densification and density data for nominally pure ZrB2 and (Zr,W)B2 specimens.  Designation  ZrB2 ZrB2 + 4 mol% W ZrB2 + 6 mol% W ZrB2 + 8 mol% W  B4C area fraction (%)  Theoretical density (g/cm3)  Bulk density (g/cm3)  Relative density (%)  4.8 1.6 1.1 0.8  5.92 6.31 6.47 6.62  5.87 6.08 6.22 6.41  99.2 98.7 98.7 99.4  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  223  Fig. 1. SEM images of the fracture surfaces of nominally pure ZrB2 (a) and ZrB2 containing 4 mol% W (b), 6 mol% W (c), or 8 mol% W (d) after oxidation at 900 °C for 5 min. Note that the images in this figure have a different magnification than those in Figs. 2 and 3.  consists of crystalline oxides containing Zr, and/or W, plus an amorphous phase that is mainly B2O3. According to the ZrO2- WO3 phase diagram [37], the presence of small (less than 5 mol%) concentrations of W should lead to the formation of Wdoped ZrO2. Higher amounts of W result in formation of ZrO2 and WO3 (up to 1105 °C) or ZrO2 and ZrW2O8 (from 1105 to 1257 °C).  ðZr1\\x00xWx ÞB2ðsÞ þ ðx þ 5ÞO2ðgÞ ! ð1 \\x00 xÞZrO2ðsÞ þ xWO3ðsÞ þ B2O3ðlÞ  ð2Þ  Figs. 1 and 2 show cross sectional SEM images of the oxide scales of nominally pure ZrB2 and ZrB2 containing 4, 6, or 8 mol% W after oxidation at 900 and 1300 °C, respectively. The oxide scales on these specimens consisted of two distinct layers: (1) a dense outer glassy layer, and (2) a porous inner layer. Due to low sensitivity of EDS to light elements, quantification of the boron content in the glassy phase was not possible. However, EDS results indicated that the matrix of the glassy phase contained O and W along  with a small amount of Zr, presumably all dissolved in B2O3. According to the ZrO2-B2O3 [38] and WO3-B2O3 [39] phase diagrams, approximately 10 mol% WO3 can dissolve into B2O3 at 900 °C, while ZrO2 is not soluble in pure B2O3 at that temperature. At 1300 °C, the solubility of ZrO2 and WO3 in B2O3 increases to 2 and 45 mol%, respectively. Immiscibility of ZrO2 and B2O3 along with the large volume expansion (\\x18300% based on density calculations) associated with oxidation of ZrB2 to ZrO2 and B2O3 are believed to result in the formation of two-layer scales [4]. After oxidation at 1600 °C (Fig. 3), the scales formed on all of the specimens consisted of a porous layer with only some residual glassy phase observed between the oxide particles. The thicknesses of the glassy and porous layers after oxidation of nominally pure ZrB2 and (Zr,W)B2 containing 4, 6, an 8 mol% W at temperatures from 800 °C to 1600 °C are summarized in Fig. 4. Between 800 and 1000 °C, the thickness of the glassy layers increased from \\x183 to \\x185 lm. The differences in the thickness of the glassy layers for nominally pure ZrB2 and (Zr,W)B2 specimens were less than 1 lm in this temperature range. Hence, the addition of W  Fig. 2. SEM images of the fracture surfaces of nominally pure ZrB2 (a) and ZrB2 containing 4 mol% W (b), 6 mol% W (c), or 8 mol% W (d) after oxidation at 1300 °C for 5 min.  \\x0c', '224  M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  Fig. 3. SEM images of the fracture surfaces of nominally pure ZrB2 (a) and ZrB2 containing 4 mol% W (b), 6 mol% W (c), or 8 mol% W (d) after oxidation at 1600 °C for 5 min.  did not have a significant effect on the thickness of the glassy layer after oxidation at temperatures below 1000 °C. At temperatures ranging from 1000 to 1200 °C, the increase in the thickness of the glassy layers was less for nominally pure ZrB2 than (Zr,W)B2 specimens, presumably due to lower evaporation of B2O3 from the (Zr,W)B2 compositions. The highest thickness of the glassy layer for nominally pure ZrB2 was 5 lm at 1200 °C, while the glassy layer was 8 lm thick for (Zr,W)B2 containing 4 mol% W, 11 lm for 6 mol% W, and 7 lm for 8 mol% W at 1300 °C. Even if it has no other effects, the dissolution of WO3 into B2O3 at elevated temperatures should lower the activity of B2O3 and, consequently, lower the evaporation of B2O3 from the liquid/glassy phase. The thickness of the glassy layers decreased after oxidation at temperatures above 1200 °C or 1300 °C for nominally pure ZrB2 and (Zr,W)B2 specimens, respectively. However, the differences between the thicknesses of the glassy layers for nominally pure ZrB2 and (Zr,W)B2 specimens were still significant (up to 14 lm) with W additions promoting the formation of thicker glassy layers compared to nominally pure ZrB2. No glassy layer was observed on the surface of nominally pure ZrB2 at 1500 °C or (Zr,W)B2 specimens at 1600 °C. After oxidation at temperatures up to 1400 °C, the thicknesses of the porous layers were almost the same for all of the compositions. Oxidation at 1600 °C produced a significantly thicker porous (\\x1885 lm) oxide layer on nominally pure ZrB2 compared to the containing W (\\x1830 lm). No considerable differences specimens were observed between the porous layer thicknesses for specimens with different amounts of W. Based on these measurements, the addition of W shifted the temperature at which the protective liquid/glassy layer evaporated and resulted in higher oxidation resistance of ZrB2 specimens with W additions compared to nominally pure ZrB2. To compare the evaporation of B2O3 from the ZrB2 and (Zr,W)B2 specimens, Eq. (1) was used to predict the thickness of the glassy scales. For this calculation, oxidation was assumed to be stoichiometric so that equal molar amounts of ZrO2 and B2O3 were formed. Building on that assumption, the glassy scale thickness  was estimated from the thickness of the porous oxide scale assuming a relative density of 98% up to 1200 °C (i.e., below the monoclinic to tetragonal phase transition temperature) and 93% at higher temperatures. The results of the predictions are shown in Fig. 5. Next, the amount of the glassy scale that evaporated during oxidation was estimated by subtracting the measured thickness from the estimated thickness (Fig. 6). Based on this analysis, the onset of evaporation increased from \\x18900 °C for nominally pure to \\x181300 °C for ZrB2 + 4 mol% W. The evaporation of B2O3 ZrB2 from ZrB2 + 4 mol% W was significantly lower than for ZrB2 for temperatures up to \\x181500 °C. Hence, the presence of W in ZrB2 suppressed evaporation of the glassy scale at temperatures below 1500 °C. Fig. 7 shows the weight gain for nominally pure ZrB2 and (Zr,W)B2 specimens containing 4, 6, or 8 mol% W after oxidation from 800 to 1600 °C in flowing air. The W content did not have a significant effect on the weight gain of ZrB2 after oxidation at temperatures below 1000 °C. The weight gains were \\x181 mg/cm2 for ZrB2 and (Zr,W)B2 specimens at 1000 °C. The weight gains for (Zr,W)B2 specimens were nearly constant between 1000 °C and 1200 °C, while for nominally pure ZrB2 weight gain increased to \\x181.5 mg/cm2 at 1100 °C and then decreased to \\x181 mg/cm2 at 1300 °C. At 1500 °C, the weight gain of (Zr,W)B2 specimens increased to \\x184 mg/cm2, while it was \\x183 mg/cm2 for nominally pure ZrB2. The weight gain for (Zr,W)B2 decreased significantly from \\x184 mg/cm2 at 1500 °C to \\x182 mg/cm2 at 1600 °C, while the weight gain of nominally pure ZrB2 increased from \\x183 mg/cm2 at 1500 °C to \\x185 mg/cm2 at 1600 °C. The measured weight gain values are representative of both the weight of oxygen added to the system due to formation of ZrO2, B2O3, and WO3 for (Zr,W)B2 and weight loss due to evaporation of B2O3 and probably WO3 for (Zr,W)B2 specimens. The increase in weight gain from 800 to 1000 °C is the result of an increase in oxidation along with insignificant evaporation of B2O3. The constant weight gain for the (Zr,W)B2 specimens from 1000 to 1200 °C indicates almost equal weights of oxygen added and B2O3 lost from the specimens. Between 1300 and 1500 °C, the evaporation of B2O3 from nominally pure  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  225  Fig. 5. Predicted thicknesses of the glassy scales without evaporation estimated from the thickness of the porous oxide scales, comparing nominally pure ZrB2 to ZrB2 containing 4 mol% W after oxidation for 5 min.  Fig. 4. (a) Glassy layer and (b) porous layer scale thicknesses as a function of oxidation temperature comparing nominally pure ZrB2 to ZrB2 containing 4, 6, or 8 mol% W after oxidation for 5 min. Note the difference in the y-axis scale between the two plots.  increases from \\x1865% to 100% of ZrB2 the glassy layer thickness, while it increases from 0% to 95% for (Zr,W)B2 specimens (see Fig. 6). On the other hand, the thickness values for the porous oxide layers (Fig. 4) are almost the same over this range of temperatures. This results in higher weight gain for (Zr,W)B2 specimens compared to nominally pure ZrB2. At 1600 °C, the higher weight gain for nominally pure ZrB2 is due to its more rapid oxidation due to the nearly complete evaporation of B2O3. The thicknesses of the porous layers are lower for (Zr,W)B2 specimens at 1600 °C, indicating less oxidation of W-containing compositions. According to the WO3-B2O3 phase diagram [39], the presence of small concentrations (<10 mol%) of WO3 with B2O3 results in the formation of a liquid phase above 800 °C. Dissolution of WO3 into the liquid phase should decrease the activity of B2O3 and, consequently, reduce its vapor pressure and evaporation rate. Further, the presence of W in the glass may also affect the B coordination state in the B2O3 network, which could further decrease the vapor pressure of B2O3 for W-containing compositions [40,41]. To inves Fig. 6. The percentage of the glassy scale lost to evaporation estimated by subtracting the measured thickness values for the glassy scales from the estimated thickness without evaporation, comparing nominally pure ZrB2 to ZrB2 containing 4 mol% W after oxidation for 5 min.  tigate the effect of WO3 additions on evaporation of B2O3, thermogravimetric analysis (TGA) was performed on pure B2O3 and borate glasses containing 4, 6 and 8 mol% WO3 at temperatures ranging from 25 to 1500 °C. The total weight loss curves as function of temperature are shown in Fig. 8. To minimize the influence of different levels of hydration among the different compositions, all of the curves were normalized to the weight at 500 °C, which is above the decomposition temperature for boric acid, but below the temperature at which significant evaporation of B2O3 begins. The weight loss for pure B2O3 was less than about 0.5% between 500 and 1000 °C. Weight loss increased linearly from \\x180.5% at 1000 °C to \\x185% at 1400 °C, and showed a sharp increase above 1400 °C. In comparison, the B2O3-WO3 glasses had weight losses of less than 0.6% up to 1300 °C. Above 1300 °C, they showed a sharp increase in weight loss. The glasses with higher amounts of losses between 1300 and 1500 °C, WO3 showed higher weight probably indicating the volatility of WO3 in this temperature range.  \\x0c', '226  M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  Fig. 7. Weight gain as a function of oxidation temperature for nominally pure ZrB2 and ZrB2 containing 4, 6, or 8 mol% W after oxidation at 800-1600 °C for 5 min in flowing air.  (Fig. 7) did not show any significant differences between ZrB2 and (Zr,W)B2 up to \\x181300 °C, despite the increased stability of the outer glassy layer for W-containing compositions. Between 1300 and 1500 °C, (Zr,W)B2 compositions showed higher mass gains than ZrB2 (Fig. 7), indicating the higher stability of the outer glassy layer. Hence, one impact of W additions to ZrB2 is to stabilize the outer glassy layer. Based on these results, as shown in Fig. 9, the addition of W to ZrB2 appears to extend the low temperature regime of oxidation behavior, where a layer of W-containing liquid B2O3 is stable on the surface of the oxidizing (Zr,W)B2. The thickness of the porous oxide layers show a sharp increase above 1400 °C due to the increased evaporation rate of B2O3 at these temperatures (Fig. 8) that results in a lower thickness of the protective liquid/glassy layers. However, the difference between the porous scales on ZrB2 and (Zr,W)B2 specimens were not significant below 1500 °C. As shown by TGA weight loss curves, at 1500 °C, the weight loss of pure B2O3 (\\x1820%) was significantly higher than B2O3-WO3 melts (<15%). The increased stability of the protective liquid/glassy layers formed on the surface of (Zr,W)B2 specimens, to higher temperatures compared to ZrB2, resulted in significantly lower weight gain and scale thickness for (Zr,W)B2 specimens compared to nominally pure ZrB2 after oxidation between 1500 and 1600 °C. Above 1500 °C for nominally pure ZrB2 and 1600 °C for (Zr,W)B2 specimens, no protective glassy layer was present on the surface of the specimens and oxidation was more severe compared to earlier regimes. However, B2O3 continues to form at these temperatures and the pores in the porous oxide layers are filled with borate melt. In the highest temperature regime, where the pores are depleted of B2O3, oxidation behavior is controlled by the flow (i.e., Knudsen diffusion as described by Parthasarathy et al. [12,13]) of molecular oxygen through the porous ZrO2 scale. Evaluating the reliability of the weight gain measurements versus the weight loss calculated from the TGA results, as a reference for the evaporation of B2O3 during oxidation, could be useful for future studies. The thickness of the glassy scales for ZrB2 and ZrB2 + 4 mol% W were calculated using the porous scale thickness measurements with evaporation calculated from weight gain (Fig. 7) and the TGA data (Fig. 8). The results are shown in Fig. 10. Again, the porosity of the porous scales were assumed to  Fig. 8. Weight loss as a function of temperature for B2O3 and borate containing 4, 6 and 8 mol% WO3 obtained by thermogravimetric analysis.  glasses  Interestingly, B2O3 containing 4 mol% WO3 exhibited the lowest weight loss of any of the compositions. Based on these observations, the presence of WO3 reduces the volatility of the borate glass up to about 1300 °C. Combined with the scale thickness and mass change results, the TGA weight loss measurements reveal several similarities and differences between the oxidation behavior of nominally pure ZrB2 and (Zr,W)B2. In the low temperature regime, below about 1000 °C, no significant differences were observed. Both ZrB2 and (Zr,W)B2 have similar scale thickness and mass gain values. Above about 1000 °C, B2O3 begins to evaporate. As indicated by TGA weight loss, pure B2O3 becomes volatile at these temperatures, which leads to the observed decrease in the thickness of the outer glassy layer on ZrB2 as shown in Fig. 4. In contrast, TGA showed that W-containing B2O3 exhibited lower weight losses, with almost no loss of mass below 1300 °C. The presence of W in the outer glassy scale reduced its volatility, which resulted in thicker glassy layers on (Zr,W)B2 up to \\x181500 °C (Fig. 4). The overall mass change  Fig. 9. TGA results showing lower weight loss for B2O3 + 4 mol% WO3 compared to B2O3 that results in lower glassy thicknesses for ZrB2 + 4 mol% W compared to ZrB2 after oxidation at 800-1600 °C for 5 min. The onset of the second oxidation regime shifts toward higher temperatures for ZrB2 + 4 mol% W.  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  227  nominally pure ZrB2 and (Zr,W)B2 compositions at temperatures between 800 and 1000 °C. However, the glassy layer thicknesses and weight gains were higher for (Zr,W)B2 specimens after oxidation from 1100 to 1300 °C. It was concluded that dissolution of W into the B2O3 liquid phase increased the stability of the protective liquid layer by reducing the volatility of B2O3 from the liquid phase, resulting in a shift in the onset of the second oxidation regime toward higher temperatures for (Zr,W)B2 specimens. This assumption was confirmed by TGA analysis of B2O3 and B2O3-WO3 glasses. Above 1500 °C, the outer glassy layer was removed from the surface of nominally pure ZrB2 by evaporation, while a thin (up to \\x183 lm) glassy layer was still covering the (Zr,W)B2 specimens. Further, the (Zr,W)B2 compositions showed lower weight gains and had thinner oxidation scales compared to nominally pure ZrB2 at 1600 °C. The addition of W into B2O3 increased the stability of the protective liquid/glassy layer and resulted in higher oxidation resistance for (Zr,W)B2 compared to nominally pure ZrB2.  Acknowledgments  This work was supported as part of the National Hypersonic Science Center for Materials and Structures (Grant FA9550-09-10477) with Dr. Ali Sayir (AFOSR) and Dr. Anthony Calomino (NASA) as program managers. The authors wish to thank project principal investigator Dr. David Marshall of Teledyne Scientific and Imaging for his support and guidance.  References  of  in dry  air,  J.  [8]  Zirconium and  ZrB2-based ceramics  [1] S.Q. 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Henderson, Convection patterns in liquid oxide films on ZrB2-SiC composites oxidized at a high temperature, J. Am. Ceram. Soc. 90 (2007) 2863-2867.  Fig. 10. Thicknesses of the glassy scales calculated from weight gains results for ZrB2 and ZrB2 + 4 mol% W after oxidation for 5 min.  and TG  be 98% up to 1200 °C and 93% at higher temperatures. The weight losses obtained by TGA were multiplied by 6 to compensate for the differences between the air flow rates in the isothermal oxidation and TGA tests. The thickness values calculated from the TGA results showed to be a good fit to the experimental results at most of the oxidation temperatures. The thickness values calculated from the weight gain results did not fit the experimental data. The calculated results indicate that weight loss obtained by TGA is more reliable than weight loss measurements made on bulk specimens that were oxidized and then cooled with respect to the evaporation of B2O3 during oxidation.  4. Conclusion  The oxidation behavior of nominally pure ZrB2 and (Zr,W)B2 ceramics with 4, 6, or 8 mol% W was studied at temperatures ranging from 800 to 1600 °C. Oxidation in this temperature range resulted in the formation of a two-layer scale: (1) an outer glassy layer containing B2O3 with dissolved W for (Zr,W)B2 compositions; and (2) a porous layer composed of oxide particles containing Zr and W for (Zr,W)B2 compositions. Based on scale thickness and weight gain measurements, two regimes of oxidation behavior were observed. The first stage was below 1100 °C for nominally pure ZrB2 and 1300 °C for (Zr,W)B2 specimens. No significant differences were measured for weight loss or scale thickness between  \\x0c', '228  M. Kazemzadeh Dehdashti et al. / Corrosion Science 80 (2014) 221-228  [20]  I. Akin, F. Cinar Sahin, O. Yucel, G. Goller, Oxidation behavior of zirconium diboride-silicon carbide composites, in: 34th International Conference on Advanced Ceramics and Composites, Daytona Beach, FL, 2010, pp. 105-111. [21] S.N. Karlsdottir, J.W. Halloran, Oxidation of ZrB2-SiC: influence of SiC content on solid and liquid oxide phase formation, J. Am. Ceram. Soc. 92 (2009) 481- 486. [22] M. Mallik, K.K. Ray, R. Mitra, Oxidation behavior of hot pressed ZrB2-SiC and HfB2-SiC composites, J. Eur. Ceram. Soc. 31 (2011) 199-215. [23] E.J. Opila, M.C. Halbig, Oxidation of ZrB2-SiC, in: 25th Annual Conference on Composites, Advanced Ceramics, Materials and Structures, Cocoa Beach, FL, 2001, pp. 221-228. [24] S. Guo, T. Mizuguchi, M. Ikegami, Y. Kagawa, Oxidation behavior of ZrB2- MoSi2-SiC composites in air at 1500 °C, Ceram. Int. 37 (2011) 585-591. [25] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, Characterization of zirconium diboride for thermal protection systems, in: 8th Conference and Exhibition of the European Ceramic Society, Istanbul, Turkey, 2004, pp. 493- 496. [26] D. Sciti, M. Brach, A. Bellosi, Oxidation behavior of a pressureless ZrB2-MoSi2 ceramic composite, J. Mater. Res. 20 (2005) 922-930. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, High-temperature chemistry and oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3, and TaSi2, J. Am. Ceram. Soc. 91 (2008) 2250-2257. [28] D. Sciti, V. Medri, L. Silvestroni, Oxidation behaviour of HfB2-15 vol.% TaSi2 at low, intermediate and high temperatures, Scripta Mater. 63 (2010) 601-604. [29] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000 °C + hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887-5904.  sintered  [27]  [33]  ultra-high  temperature  [30] E.J. Opila, M.C. Halbig, Oxidation of ZrB2-based ceramics, Ceram. Eng. Sci. Proc. 22 (2001) 221-228. [31] S.R. Levine, E.J. Opila, Tantalum Addition to Zirconium Diboride for Improved Oxidation Resistance, in: NASA/TM-2003-212483, 2003. [32] F. Peng, Y. Berta, R.F. Speyer, Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride, J. Mater. Sci. 24 (2009) 1855-1867. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, S. Dallek, Oxidation of ZrB2 ceramics modified with SiC and group IV-VI transition metal diborides, Electrochem. Soc. Proc. 12 (2001) 144-158. [34] P. Hu, X.H. Zhang, J.C. Han, X.G. Luo, S.Y. Du, Effect of various additives on the oxidation behavior of ZrB2-based ultra-high-temperature ceramics at 1800 °C, J. Am. Ceram. Soc. 93 (2010) 345-349. [35] X.H. Zhang, P. Hu, J.C. Han, L. Xu, S.H. Meng, The addition of lanthanum hexaboride to zirconium diboride for improved oxidation resistance, Scripta Mater. 57 (2007) 1036-1039. [36] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Oxidation of zirconium diboride with tungsten carbide additions, J. Am. Ceram. Soc. 94 (2011) 1198-1205. [37] L.L.Y. Chang, M.G. Scroger, B. Phillips, Condensed phase relations in the systems ZrO2-WO2-WO3 and HfO2-WO2-WO3, J. Am. Ceram. Soc. 50 (1967) 211-215. [38] S.N. Karlsdottir, J.W. Halloran, A.N. Grundy, Zirconia transport by liquid convection during oxidation of zirconium diboride-silicon carbide, J. Am. Ceram. Soc. 91 (2008) 272-277. [39] E.M. Levin, The system WO3-B2O3, J. Am. Ceram. Soc. 48 (1965) 491-492. [40] G. Pal Singh, D.P. Singh, Structural and optical properties of WO3-PbO-B2O3 glass-ceramic, J. Phys. Chem. Solids 73 (2012) 540-544. [41] W. Vogel, Glass Chemistry, second ed., Springer-Verlag, New York, 1994.  \\x0c']"
},{
  "_id": 64,
  "PDF": "Effects of the second phase on the microstructure and ablation resistance o HfC coating on C-C composites.pdf",
  "Text": "[\"Surface & Coatings Technology 344 (2018) 250-258  Contents lists available at ScienceDirect  Surface & Coatings Technology  jou rna l homepage : www .e lsev ie r .com / loca te /su r fcoa t  Effects of the second phase on the microstructure and ablation resistance of HfC coating on C/C composites  T  Jincui Rena, Yulei Zhanga,⁎, Yanqin Fua, Pengfei Zhanga, Song Tianb, Leilei Zhanga,⁎  a State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi'an 710072, China b School of Materials Science and Engineering, Chongqing Jiaotong University, Chongqing 400074, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: HfC coating Supersonic atmospheric plasma spraying Second phase Bonding strength Ablation resistance  To improve the ablation resistance of HfC coating on carbon/carbon composites prepared by supersonic atmospheric plasma spraying, SiC, TaC and ZrC ceramics were added into HfC coating as the second phase. Effects of the second phase on the microstructure, interface bonding strength and ablation resistance of HfC coating were researched. The structure of HfC-SiC coating was denser and smoother than that of HfC-TaC and HfC-ZrC coatings. The dense structure, high interface bonding strength and formation of continuous Hf-Si-O glassy film resulted in the better ablation resistance of HfC-SiC coating. During ablation, the HfC-TaC coating was oxidized to a stable Hf6Ta2O17 glassy film, which could restrain the oxygen penetration. A porous HfO2-ZrO2 layer was generated on the HfC-ZrC coating after ablation, some HfO2 and ZrO2 particles may be flown away by the highspeed flame. So, the ablation resistance of HfC-ZrC coating was worse than the other two coatings.  1.  Introduction  Carbon/carbon (C/C) composites are noted for low density, low coefficient of thermal expansion (CTE) (1.0 × 10−6 °C−1), high thermal conductivity and good high temperature mechanical properties [1-4]. These good properties make them to be promising high temperature structural materials in aerospace industry [5,6]. However, C/C composites are easily oxidized above 450 °C, resulting in the degradation of their mechanical properties [7]. Applying a ceramic coating is an effective method to protect C/C composites from oxidation [4-7]. Due to the good oxidation resistance and relatively low CTE of SiC ceramic, SiC-based coatings have showed outstanding oxidation and thermal shock resistance below 1700 °C [8,9]. While, the melting point of SiO2 is 1600-1700 °C, in ultra-high temperature oxidizing environment with high-speed gas flow washout, the SiC-based coatings would be oxidized, melted and blown away. Refractory metal carbides and borides are suitable coating materials to improve the ablation resistance of C/C composites due to their high melting point, high hardness and high chemical stability [10-15]. It has been reported that the vapor pressure of carbides is lower than that of borides [16], so refractory metal carbides are more suitable to be used as the ablation resistance coating materials. Among all of the refractory metal carbides, HfC is characterized by the highest melting temperature (3959 °C) [17], high mechanical properties, good resistance to chemical attack and ablation, and high  temperature phase stability [12,18]. In addition, the melting point of HfO2 is 2810 °C, which is higher than that of ZrO2 (2667 °C) and Ta2O5 (1800 °C). So, from the perspective of thermal stability, HfC is a promising coating material to protect C/C composites from ablation in severe ablation environment. During ablation, HfC coating will be oxidized to porous HfO2 layer. The pores in the oxide coating will provide penetration channels for oxygen. In addition, the chalked oxide coating will easily peel off from C/C substrate. Introducing carbides with lower melting point as the second phase could improve the oxidation resistance of HfC ceramic [19,20]. Xiong et al. [21] synthesized HfC/ZrC biphasic coating on C/C composites by chemical vapor deposition (CVD), results showed that ZrC could improve the ablation resistance of HfC coating. Wang et al. [22] synthesized Hf(Ta)C coating by CVD, which indicated that a dense and continuous Hf6Ta6O17 glassy layer was generated during ablation, which could inhibit the oxygen infiltration and improve the ablation resistance of the coating. Yang et al. [20] prepared HfC-SiC coating by supersonic atmospheric plasma spraying (SAPS), results indicated that a Hf-Si-O compound oxide layer was formed during ablation, which could restrain the oxygen penetration. So, introducing carbide with lower melting point as the second phase could improve the ablation resistance of HfC coating. CVD has been applied to fabricate HfC single-phase coating and HfC-based composite coating [6,12,21-23]. While, the deposition process is usually time-consuming and high-cost. SAPS technology is characterized by high deposition rate and low cost, which has been  ⁎ Corresponding authors. E-mail address: zhangyulei@nwpu.edu.cn (Y. Zhang).  https://doi.org/10.1016/j.surfcoat.2018.03.023 Received 2 November 2017; Received in revised form 24 January 2018; Accepted 11 March 2018  Available online 12 March 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.  \\x0c\", \"J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  Fig. 1. XRD patterns of the as-prepared HfC-based biphasic coatings: (a) HfC-SiC coating; (b) HfC-TaC coating; (c) HfC-ZrC coating.  widely applied to synthesize coatings. During SAPS process, the temperature of plasma arc is very high (above 10,000 °C) and the velocity of spraying particle is supersonic [21,24-26]. The composition, microstructure and thickness of the coatings could be designed by adjusting the spraying parameters. So, SAPS is a promising method to prepare refractory metal carbides coatings. In this work, HfC-SiC, HfC-TaC and HfC-ZrC coatings were prepared on SiC-coated C/C composites by SAPS. Effects of the second phase on the microstructure, bonding strength and ablation resistance of HfC coating were investigated.  2. Experimental procedure  2.1. Synthesize of HfC-based coatings on SiC-coated C/C composites  Two-dimensional (2D) C/C composites with a density of 1.65 g/cm3 were used as substrates. The bulk C/C composites were cut into two kinds of small specimens: 10 × 10 × 10 mm3 for morphology characterization and bonding strength test; φ30 × 10 mm3 for ablation test. The specimens were polished using 400 grit SiC paper, ultrasonically cleaned by distilled water and dried at 85 °C for 2 h. SiC inner coating was prepared on C/C substrates by pack cementation. The mixture powders composed of 60-80 wt% Si, 15-25 wt% C and 5-10 wt% Al2O3 were used as raw materials. C/C specimens were embedded in the mixture powders and then heated to 1800-2100 °C for 2 h. The detailed preparation process has been reported in Ref. [27]. HfC-based coatings were prepared on SiC-coated C/C specimens by SAPS using a HEPJ-II SAPS equipment (Beijing Armored Force Engineering College). Commercial HfC, SiC, TaC and ZrC powders were used as raw materials (400 mesh; 99 wt% purity; Jinzhou Metal  Material Research Institute, China). HfC-15 vol% SiC, HfC-15 vol% TaC and HfC-15 vol% ZrC powders were mixed by a blender for 3 h. To improve the flowability of the powders, a slurry consisting of 48 wt% distilled water, 2 wt% polymeric binder and 50 wt% mixture powders was agglomerated and pelleted using a spray drier, then sifted by 200 mesh sieve. During spraying, the specimens were placed perpendicularly to the plasma torch. The spraying distance was 100 mm. Ar was used as the primary gas and carrier gas. H2 was used as the secondary gas. The flow rate of Ar and H2 was 80 and 5 L/min, respectively. The spraying current and voltage was 400 A and 125 V, respectively. The feed rate of the spraying powders was 25 g/min.  2.2. Phase composition and microstructure characterization of the coatings  The phase composition and morphology of the coatings were analyzed by X-ray diffraction (XRD, X'Pert Pro, PANalytical, Almelo, the Netherlands) and field emission scanning electron microscopy (SEM, TESCAN VEGA3, Czech), respectively. The phase distribution of the coatings were analyzed through back scattered electron (BSE) image of the coatings.  2.3. Bonding strength and ablation tests for the coatings  Scratch test was performed using WS-2015multi-functional tester to measure the interface bonding strength between SiC coating and HfCbased coatings. A load increasing from 0 to 20 N was applied on the coating surface with a sliding distance of 5 mm. The acoustic emission (AE) signal-load curve was recorded by computer. The final bonding strength was the average value of five samples. The ablation test was performed under oxyacetylene flame. During  251  \\x0c\", 'J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  )a(  (c)   (e)   fullly molten area   insufficiently molten area   )b(  (d)   (f)   Fig. 2. Surface SEM images of the HfC-based biphasic coatings: (a)(b) HfC-SiC coating; (c)(d) HfC-TaC coating; (e)(f) HfC-ZrC coating.  ablation, the flow rate of oxygen and acetylene was 0.244 and 0.167 L/ s, and the pressure of them was 0.4 and 0.095 MPa. The ablation angle was 90°. The ablation distance was 10 mm. The surface temperature of the coated samples in central region was recorded by an infrared thermometer. After ablation for 120 s, the samples were naturally cooled to room temperature. The mass (Rm) and linear (Rl) ablation rates of the samples were calculated based on the following equations:  R  m  = Δm t /  R  l  Δ / d t  =  (1)  (2)  Δm is the mass loss of the samples after ablation; Δd is the decreased thickness of the samples in center region after ablation; t is the ablation time. The final ablation rates were the average values of three samples.  3. Results and discussion  3.1. Phase composition and morphology of  the HfC-based coatings  Fig. 1 shows the XRD patterns of the HfC-based coatings. As seen in Fig. 1(a), the HfC-SiC coating consists of HfC, SiC, HfO2 and SiO2 phases. The spraying process was performed in air environment, so some HfC and SiC powders were oxidized to HfO2 and SiO2. The content of SiC in the spraying powders is low (15 vol%), so SiC and SiO2 peaks are weak. The HfC-TaC coating is composed of HfC, TaC, HfO2 and Ta2O5 phases (Fig. 1(b)). During spraying, the oxides of HfC and TaC (HfO2, Ta2O5) were generated. TaC and Ta2O5 peaks are weak because  of the low content of TaC in the raw materials (15 vol%). Fig. 1(c) reveals the XRD pattern of the HfC-ZrC coating, from which HfC, ZrC and their oxides (HfO2/ZrO2) peaks can be detected. The diffraction peaks of HfO2 and ZrO2 are very close and difficult to distinguish [21]. Fig. 2 illustrates the surface morphologies of the HfC-based coatings. The three coatings are rough, consisting of fully molten area and insufficiently molten area. In spraying process, the spraying powders quickly melted and solidified on the sample surface to form splats. The fully molten area is smooth and dense without void (Fig. 2(b)). Some powders did not completely melt, leading to the formation of voids in insufficiently molten area (Fig. 2(d) and (f)). The melting point of SiC and its oxide is lower than that of TaC, ZrC and their oxides. So, HfC-SiC coating is mainly composed of fully molten area. The structure of HfCSiC coating is denser and flatter than that of HfC-TaC and HfC-ZrC coatings. As seen in Fig. 2(d), some voids are distributed on the HfCTaC coating, the insufficiently molten area is larger due to the higher melting point of TaC and its oxide. As shown in Fig. 2(f), the HfC-ZrC coatings is loose, the unmelted powders stack up and form many voids in the coating. The cross-section BSE images of the coatings are displayed in Fig. 3. All of the three coatings are composed of two layers: inner SiC coating and outer HfC-based coating. The interface between SiC coating and C/ C substrate is jagged, indicating a good bonding between them. The CTE of SiC is relatively low (4.5 × 10−6 °C−1), so the SiC coating could relieve the CTE mismatch between C/C substrate and the outer HfCbased coatings. A distinct interface can be found between inner and  252  \\x0c', 'J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  (a)   outer coating   (b)   C/C   inner coating   (c)   interface gap   uneven area   Fig. 3. Cross-section SEM images of the HfC-based biphasic coatings: (a) HfC-SiC coating; (b) HfC-TaC coating; (c) HfC-ZrC coating.  Fig. 4. Variation of acoustic emission signal with sliding load during scratch test: (a) HfC-SiC-coated sample; (b) HfC-TaC-coated sample; (c) HfC-ZrC-coated sample.  253  \\x0c', 'Surface & Coatings Technology 344 (2018) 250-258  coatings. The bonding strength between SiC coating and HfC-based coatings was measured by scratch test. Fig. 4 shows the AE-load curve during the scratch test. Appearance of the first obvious AE signal indicates the spalling of outer coating from the inner coating [28]. Result shows that the bonding strength between SiC coating and HfC-SiC, HfCTaC and HfC-ZrC coatings is 14.5 ± 1.24, 10.8 ± 0.86 and 9 ± 0.57 N, respectively. The HfC-SiC coating is uniform and compact, no void or gap is at the interface between the two coatings. So, the bonding strength between SiC coating and HfC-SiC coatings is the highest. In contrast, voids and interface gap exist in the HfC-ZrC coating. Therefore, the bonding strength between SiC coating and HfCZrC coatings is the lowest. Table 1 shows the ablation performance of the coatings. After ablation for 120 s, the Rl of the HfC-SiC, HfC-TaC and HfC-ZrC-coated samples is −0.67, −0.58 and − 0.42 μm/s, respectively. The Rl of the samples are all negative, resulting from the volume expansion of the coatings. The oxidation of carbides and phase transition of oxides (HfO2 and ZrO2) during quick cooling will lead to a volume expansion [19]. As seen in Table 1, the Rm of the HfC-SiC, HfC-TaC and HfC-ZrC-coated samples is −0.15, 0.68 and 1.23 mg/s, respectively. The weight gain of HfC-SiC-coated sample is because of the coating oxidation. In contrast, the weights of HfC-TaC and HfC-ZrC-coated samples are both reduced after ablation due to the mechanical denudation of oxide layer by gas flow. The Rl and Rm of the HfC-SiC-coated sample are lower than those of HfC-TaC and HfC-ZrC-coated samples, indicating that the HfC-SiC coating has the best ablation performance among the three coatings. During ablation, central region suffered the highest temperature and the most serious mechanical denudation, leading to the most serious ablation in central region. Fig. 5 exhibits the surface ablation temperature curves of the coated samples in central region. The three ablation temperature curves show the same changing trend. In the initial 5 s, the surface temperature of the samples rapidly increased to 1500-2000 °C. Then, the temperature slowly rose to 1800-2300 °C and maintained stable. The highest surface temperature of HfC-SiC, HfCTaC and HfC-ZrC coated C/C sample is 1855, 2000 and 2290 °C, respectively. During ablation, HfC-SiC coating was oxidized to HfO2 and SiO2. The melting point of SiO2 (about 1700 °C) is lower than ablation temperature. The melt and volatilization of SiO2 could absorb much heat, so the surface temperature of HfC-SiC coated sample is the lowest. The melting temperature of Ta2O5 is about 1800 °C, some Ta2O5 in the coating would melt and volatilize, and some other Ta2O5 would form solid solution with HfO2. So, the surface temperature of HfC-TaC coated sample is slight higher than that of HfC-SiC coated sample. The oxides of HfC-ZrC coating have higher melting point (HfO2, 2810 °C; ZrO2, 2667 °C), only some oxides could melt during ablation. Therefore, the surface temperature of HfC-ZrC coated sample is the highest. At higher temperature, the oxidation and ablation rate of the coating is higher, resulting in the worse ablation resistance. Fig. 6 displays the macrographs of the samples after ablation. Three  J. Ren et al.  Table 1 Ablation property of  the coated C/C samples.  HfC-based coatings  Linear ablation rates [μm/s]  Mass ablation rates [mg/s]  HfC-SiC coating HfC-TaC coating HfC-ZrC coating  −0.67 ± 0.12 −0.58 ± 0.07 −0.42 ± 0.08  −0.15 ± 0.03 0.68 ± 0.11 1.23 ± 0.18  Fig. 5. Surface temperature curves of  the coated C/C samples  in central  region during  ablation.  outer coatings, suggesting that the bonding between them is mechanical. The thickness of SiC coating and outer HfC-based coatings is about 50-100 and 100-200 μm, respectively. From Fig. 3(a), it can be found that a dense HfC-SiC coating is uniformly adhered to the inner SiC coating. There is no void or crack at the interface between the two coatings. The HfC-TaC coating is slightly thicker (between 125 and 150 μm), while the compactness of the coating is not very high (Fig. 3(b)). From Fig. 3(c), it can be observed that the thickness of HfCZrC coating is uneven (indicated by the yellow arrow). This is because that the powders were not completely melted during spraying, the unmolten particles stacked up and could not evenly spread out on the surface of SiC coating. Moreover, an interface gap can be found, resulting from the CTE mismatch between the SiC and HfC-ZrC coatings.  3.2. Bonding strength and ablation resistance of  the HfC-based coatings  Low bonding strength between the different coatings will lead to the outer coating cracking or spalling during ablation. So, high bonding strength is advantageous for improving the ablation resistant of the  C   A   B   HfC-SiC coating   HfC-TaC coating   HfC-ZrC coating   Fig. 6. Macrographs of  the coated C/C samples after the ablation test.  254  \\x0c', 'J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  Fig. 7. XRD patterns of  the coated C/C samples after the ablation test: (a) HfC-SiC-coated sample; (b) HfC-TaC-coated sample; (c) HfC-ZrC-coated sample.  distinct regions can be observed on the coating surface: central region (A), transition region (B) and border region (C). After ablation, all of the three samples are covered with white ablation products. The coatings maintain integrity without any peeling. The central region suffered the most severe ablation, so the quantity of ablation products in central region is the most among the three ablation regions. As seen in Fig. 6, the HfC-SiC and HfC-TaC coatings are both compact after ablation. The surface of HfC-ZrC coating is porous with some small voids. During ablation, the formed HfO2 and ZrO2 layer could not be melted, some of the solid oxide particles were washed out by gas flow, resulting in the generation of voids in the coating. XRD patterns of the samples after ablation are illustrated in Fig. 7. Because of the oxidation of HfC and SiC, monoclinic HfO2 (m-HfO2) and SiO2 were formed on HfC-SiC-coated sample (Fig. 7(a)). Some generated SiO2 would melt and evaporate, and some SiO2 may be amorphous after quick cooling, so the diffraction peaks of SiO2 are weak [20]. The diffraction peaks of HfSiO4 can also be found in the XRD pattern, resulting from the reaction between HfO2 and SiO2. XRD pattern of the HfC-TaC coating shows that the coating consists of HfO2 and Hf6Ta2O17 phases after ablation (Fig. 7(b)), indicating that the oxides of HfC and TaC formed solid solution during ablation. According to the JCPDS database, Hf6Ta2O17 is the only ternary crystalline phase in Hf-Ta-O system [29]. Some diffraction peaks of HfO2 and Hf6Ta2O17 are difficult to distinguish. From Fig. 7(c), it can be found that the ablation products of HfC-ZrC coating are HfO2 and ZrO2, whose diffraction peaks are very close and difficult to distinguish. Surface SEM morphologies of the samples in central region after ablation are exhibited in Fig. 8. As seen in Fig. 8(a), a continuous and  dense layer is formed on HfC-SiC-coated sample. The enlarged image (Fig. 8(b)) exhibits that the coating is composed of connected oxide grains. During ablation, HfC-SiC coating was oxidized to HfO2, SiO2 and HfSiO4, resulting in the generation of Hf-Si-O compound glassy layer. The evaporation of gas products (CO, SiO2) would leave small voids on the coating. A continuous glassy layer is formed on the surface of HfCTaC coating after ablation (Fig. 8(c)), which is composed of connected oxide grains. Some small voids can also be found due to the volatilization of gas products. There are some cracks on the coating (as seen in Fig. 8(d)), resulting from the CTE mismatch between inner and outer coatings. As seen in Fig. 8(e), the HfC-ZrC coating has obviously different morphology with the HfC-SiC and HfC-TaC coatings. After ablation, the oxide layer of HfC-ZrC coating is rough and discontinuous, which is composed of many small grains. The melting points of HfO2 and ZrO2 are higher than the ablation temperature, so the oxide particles were not fully melted during ablation. The oxide grains may be washed out by gas flow, resulting in the existence of voids in the oxide layer. The porous oxide layer could not restrain the oxygen infiltration, so the ablation resistance of HfC-ZrC coating is not as good as the HfCSiC and HfC-TaC coatings.  3.3. Ablation mechanism of  the HfC-based coatings  The HfC-based coatings can protect C/C substrates from oxidation by consuming oxygen and forming oxygen barrier layers. During the ablation under oxyacetylene flame [30], the following reactions may occur:  255  \\x0c', 'J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  (a)   (c)   (e)   voids   cracks   (b)   (d)   (f)   Fig. 8. Surface SEM images of the coatings in central region after ablation for 120 s: (a)(b) HfC-SiC coating; (c)(d) HfC-TaC coating; (e)(f) HfC-ZrC coating.  2HfC (s)  +  3O (g)  2  2HfO (s)  2  2CO (g)  +  →  2SiC (s)  +  3O (g)  2  2SiO (s)  2  2CO (g)  +  →  SiO (s)  2  SiO (l)  2  →  SiO (l)  2  SiO (g)  2  →  HfO (s)  2  SiO (l)  2  +  →  HfSiO (s)  4  12HfC (s)  4TaC (s)  +  +  25O (g)  2  →  Hf Ta O  2  6  17  (s)  →  Hf Ta O  6  2  17  (l)  2Hf Ta O  6  2  17  2ZrC (s)  +  3O (g)  2  2ZrO (s)  2  2CO (g)  +  →  (3)  (4)  (5)  (6)  (7)  (8)  (9)  (10)  (s)  16CO (g)  +  Fig. 9(a) reveals the ablation schematic diagram of HfC-SiC-coated sample. During ablation, HfC-SiC coating was oxidized to HfO2 and SiO2 according to the Reactions (3) and (4). A Hf-Si-O compound glassy layer was formed on the sample surface. Meanwhile, HfO2 could react with SiO2 to generate a new stable phase (HfSiO4) (Reaction (7)). During ablation, the formed Hf-Si-O glassy layer was dense and stable, which could inhibit the oxygen infiltration. The ablation temperature was higher than the melting point of SiO2, so according to Reactions (5) and (6), some generated SiO2 would melt and evaporate. The melting point of HfSiO4 is above 2900 °C, the HfSiO4 particles could have a pinning effect on the Hf-Si-O glassy layer, which could inhibit the formation and propagation of cracks. The melt and evaporation of SiO2  256  were heat absorbing process, which could consume much heat. So, the surface temperature of HfC-SiC coating was the lowest (Fig. 5). The evaporation of gaseous products (CO, SiO2) would leave small voids in the coating. Meanwhile, the glassy layer could seal some voids. So, the voids may have little effect on the ablation resistance of the coating. The ablation resistance of HfC-SiC coating was the best among the three HfC-based coatings due to the following reasons: the structure of HfCSiC coating was dense and uniform before ablation; the interface bonding strength between SiC coating and HfC-SiC coating was high; the surface ablation temperature of HfC-SiC-coated sample was low; a dense, stable and continuous Hf-Si-O glassy layer was formed during ablation. The schematic diagram of the TaC-SiC-coated sample during ablation is exhibited in Fig. 9(b). A compound oxide layer consisting HfO2 and Hf6Ta2O17 was generated on HfC-TaC coating after ablation through Reactions (3) and (8). Both HfO2 (2.09 W/m °C) and Hf6Ta2O17 (2.89 W/m °C) have low thermal conductivity [31], which could restrain the heat diffusion and then decrease the oxidation rate of the coating. The melting point of Hf6Ta2O17 is between HfO2 and Ta2O5 [22]. So, some molten Hf6Ta2O17 was formed during ablation (Reaction (9)), which could seal the voids in the porous oxide layer and decrease the infiltration rate of oxygen. The melting of Hf6Ta2O17 could consume much heat, which could decrease the surface temperature of the coated samples. In addition, the Hf6Ta2O17 has only one phase structure in the reported research [31-33], so Hf6Ta2O17 has no phase transition during  \\x0c', 'J. Ren et al.  Surface & Coatings Technology 344 (2018) 250-258  Fig. 9. Ablation schematic diagram of  the coated samples: (a) HfC-SiC-coated sample; (b) HfC-TaC-coated sample; (c) HfC-ZrC-coated sample.  cooling process. Moreover, the glassy oxide layer has high viscosity, which has good scouring resistance to gas flow. Because of the existence of more voids in the original HfC-TaC coating and lower interface bonding strength, the ablation resistance of HfC-TaC coating was worse than that of HfC-SiC coating. Fig. 9(c) illustrates the schematic diagram of ZrC-SiC-coated sample during ablation. For HfC-ZrC coating, the oxidation of carbides happened during ablation (Reactions (3) and (10)). Both the melting points of HfO2 and ZrO2 are higher, so only bits of oxides could melt. For long time ablation (120 s), the Hf-Zr-O layer would solid state sinter. On the base of the binary phase diagram of HfO2 and ZrO2, HfO2 and ZrO2 can form solid solution, which could accelerate the sintering process of oxides and improve the ablation resistance of the coating [34]. The surface temperature of HfC-ZrC-coated sample was the highest during ablation, so the oxidation rate of the coating was the fasted. There were some voids in the original HfC-ZrC coating, oxygen would infiltrate into the coating through these voids, resulting in the rapid oxidation of HfCZrC coating. Some of the HfO2 and ZrO2 particles would be washed out by the high-speed gas flow, resulting in the oxidation of inner coating. No continuous glassy layer formed on the coating surface, so the ablation resistance of HfC-ZrC coating was the worst among the three HfCbased coatings.  4. Conclusion  HfC-SiC, HfC-TaC and HfC-ZrC composites coatings were synthesized on SiC-coated C/C composites by supersonic atmospheric plasma spraying. All of the three coatings consisted of fully molten area and  257  insufficiently molten area. The surface of HfC-SiC coating was dense and uniform. In contrast, the insufficiently molten area on the HfC-TaC and HfC-ZrC coatings were more. There is an interface gap between HfC-ZrC coating and SiC coating. The interface bonding strength between SiC coating and HfC-SiC, HfC-TaC and HfC-ZrC coatings was 14.5 ± 1.24, 10.8 ± 0.86 and 9 ± 0.57 N, respectively. Due to the dense structure of HfC-SiC coating and the formation of continuous HfSi-O glassy layer, the ablation resistance of HfC-SiC coating was the best. After ablation for 120 s, the Rl and Rm of HfC-SiC-coated sample was −0.67 μm/s and − 0.15 mg/s, respectively. The HfC-TaC coating was oxidized to compound glassy layer consisting HfO2 and Hf6Ta2O17 during ablation, which could restrain the oxygen infiltration. The Rl and Rm of HfC-TaC-coated samples was −0.58 μm/s and 0.68 mg/s after ablation for 120 s. The HfC-ZrC coating was oxidized to HfO2 and ZrO2 layer after ablation, which was rough and porous. So, the ablation resistance of HfC-ZrC coating was the worst. The Rl and Rm of HfC-ZrCcoated samples was −0.42 μm/s and 1.23 mg/s after ablation for 120 s.  Acknowledgements  This work has been supported by the National Natural Science Foundation of China under Grant No. 51672223, 51521061, 51502028, the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant No.98-QZ-2014), and the Basic and Advanced Research Project of Chongqing Science and Technology Commission (Grant no. cstc2015jcyjA50011).  \\x0c', 'J. Ren et al.  References  [7]  [9]  [5]  [4]  [2]  [1]  F. Smeacetto, M. Salvo, M. Ferraris, Oxidation protective multilayer coatings for carbon-carbon composites, Carbon 40 (2002) 583-587. E.L. Corral, L.S. Walker, Improved ablation resistance of C-C composites using zirconium diboride and boron carbide, J. Eur. Ceram. Soc. 30 (2010) 2357-2364. [3] Q.G. Fu, J.Y. Jing, B.Y. Tan, R.M. Yuan, L. Zhuang, L. Li, Nanowire-toughened transition layer to improve the oxidation resistance of SiC-MoSi2-ZrB2 coating for C/C composites, Corros. Sci. 111 (2016) 259-266. T. Li, H.J. Li, X.H. Shi, J. Cheng, L. 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},{
  "_id": 65,
  "PDF": "Effects of transition metals on the oxidation behavior of ZrB2 ceramics.pdf",
  "Text": "['Corrosion Science 91 (2015) 224-231  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effects of transition metals on the oxidation behavior of ZrB2 ceramics  ⇑  M. Kazemzadeh Dehdashti, W.G. Fahrenholtz  , G.E. Hilmas  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65401, United States  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 13 July 2014 Accepted 12 November 2014 Available online 18 November 2014  Keywords:  A. Ceramic B. SEM C. Oxidation  1. Introduction  The effects of transition metal additives on the oxidation behavior of ZrB2 ceramics at 1600 °C were investigated. Additions of W, Mo, or Nb improved the oxidation resistance of ZrB2 even after evaporation of the protective B2O3 layer. The oxide layers were composed of an outer layer that was porous and had light color and an inner layer that was dense and appeared dark. Additions of Mo or Nb were more effective than W at improving the oxidation resistance of ZrB2 at 1600 °C due to a lower evaporation rate for B2O3 glasses containing Nb and Mo compared to W.  Ó 2014 Elsevier Ltd. All rights reserved.  Ultra high temperature ceramics (UHTCs) are a class of materials that includes diborides (i.e. ZrB2, HfB2), carbides (i.e. ZrC, HfC), and nitrides (i.e. ZrN, HfN). UHTCs have melting temperatures in excess of 3000 °C and the ability to withstand extreme heating environments. Diboride-based UHTCs have performance advantages compared to carbides and nitrides due to better oxidation resistance and better ability to transfer and redistribute heat (thermal conductivity >100 W/m K at 25 °C) at elevated temperatures. These characteristics make them good candidates for sharp leading edges for hypersonic aerospace vehicles, which must be capable of operating in oxidizing atmospheres at high temperatures and high flow rates, and also have high thermal conductivity to transfer heat away from the hottest areas and redistribute it to cooler areas [1,2]. Zirconium diboride is attractive for aerospace applications due (\\x183250 °C), to its high melting point low theoretical density (6.09 g/cm3), and a thermal conductivity of over 100 W/m K at room temperature [3,4]. However, oxidation at temperatures above 800 °C has limited the development of ZrB2 ceramics for aero-propulsion and hypersonic flight applications [5]. Oxidation of ZrB2 is usually assumed to be stoichiometric, resulting in measurable mass gain due to formation of B2O3 and ZrO2 according to the following reaction [6].  ZrB2ðsÞ þ 5=2 O2ðgÞ ! ZrO2ðsÞ þ B2O3ðlÞ  ð1Þ  Previous studies have recognized three different temperature regimes of oxidation for ZrB2 [7]. The low temperature regime occurs below about 1100 °C. In this regime, crystalline ZrO2 and  ⇑ Corresponding author. Tel.: +1 573 342 6343; fax: +1 573 341 6934. E-mail address: billf@mst.edu (W.G. Fahrenholtz).  http://dx.doi.org/10.1016/j.corsci.2014.11.019 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.  a continuous liquid/glassy B2O3 layer form on the surface of the unoxidized ZrB2 matrix, providing passive oxidation protection. Based on mass gain and increasing oxide scale thickness, the oxidation kinetics are parabolic in the first regime [8-12]. The second oxidation regime is between \\x181100 and \\x181400 °C. In this temperature range, the weight change represents a combination of mass gain due to formation of ZrO2 and B2O3, and mass loss due to evaporation of B2O3 [9,12,13]. Kinetics in this regime are para-linear due to the combination of the protective nature of the B2O3 layer (parabolic component) and the mass loss due to B2O3 evaporation (linear component). The third oxidation regime is above \\x181500 °C. At these temperatures, the evaporation rate of B2O3 is more significant. As a result, nearly all of the B2O3 is lost by evaporation and a porous ZrO2 scale covers the surface. In this regime, linear mass gain kinetics are observed, indicating that the oxide layers do not act as a barrier to oxygen transport [8]. The conventional approach to improving the oxidation resistance of diboride ceramics is to add Si-containing compounds such as SiC [6,14-24], MoSi2 [25-27], or TaSi2 [28,29]. The higher stability of the borosilicate glassy layer formed on the surface of Si-containing diborides, compared to the borate glassy layer that forms on nominally pure ZrB2, results in improved oxidation resistance of ZrB2-SiC ceramics compared to pure ZrB2 [6,24]. Nevertheless, Si-containing additives also result in problems, such as rupture of the protective glassy layer as a result of SiO(g) formation [30], or formation of a SiC-depleted layer [6]. These issues may result in loss of protection or damage to the underlying ceramic at temperatures above 1600 °C [31,32]. Several studies have shown that additions of transition metals (TMs), including Mo, Nb, V, Cr, Ti, or Ta, as borides or silicides can improve the oxidation resistance of ZrB2 and ZrB2-SiC ceramics [27,32-36]. However, very few studies have examined the effects of transition metal additives on the oxidation behavior of  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  225  silicon-free ZrB2 ceramics. Initial studies by Zhang et al. [13,37] reported that the additions of tungsten carbide (WC) improved the oxidation resistance of ZrB2 ceramics by formation of WO3 in the oxide scale. The studies concluded that WO3 formed during oxidation and reacted with another oxidation product, ZrO2, to form a liquid phase. Liquid phase formation promoted sintering of ZrO2 in the scale and, consequently, increased the relative density of the scale, which decreased the oxidation rate. Subsequent studies have shown that the addition of W and Nb to nominally pure ZrB2 improved its oxidation resistance by increasing the stability of the protective liquid/glassy B2O3 scale [7,38]. The effect of transition metal additives on the oxidation behavior of nominally pure ZrB2 at temperatures at which the borate glassy layer is evaporated has not been studied yet. The present study investigated the effects of W, Mo, and Nb on the oxidation behavior of ZrB2 ceramics at 1600 °C. These additives were selected based on a combination of their ability to form solid solutions with ZrB2 and miscibility with B2O3 at elevated temperatures (e.g., 1600 °C). The goals of the study were to identify the mechanisms by which the additives improved oxidation resistance after loss of the B2O3 liquid phase and to find the most effective TM for improving the oxidation resistance of ZrB2 ceramics at elevated temperatures.  after the desired oxidation time and air quenched to room temperature to minimize changes such as further oxidation or evaporation of B2O3 that may occur during cooling. Specimens that were quenched immediately after reaching the desired temperature were considered to be oxidized for ‘‘0 h’’ while other specimens were held at the oxidation temperature for up to 6 h before quenching. To study the microstructure and calculate the thicknesses of the oxide layers, fracture surfaces were observed using secondary electron scanning electron microscopy (SEM; S-4700, Hitachi, Japan) by applying accelerating voltage of 10 kV on Au- Pd coated specimens, and optical microscopy (Hirox KH-8700, Hackensack, NJ). Since B2O3 can flow from the side surfaces to the bottom surface of the specimens during oxidation, scale thickness measurements were always performed on the upper surface. The thicknesses of the oxide layers were measured using image analysis software (ImageJ, U.S. National Institutes of Health, Bethesda, MD) and are reported as an average of 20 measurements. The surface coverage was also measured from SEM images using the same software. Chemical compositions of the oxide scales were studied using energy dispersive spectroscopy (EDS; EDAX, Mahwah, NJ).  3. Results and discussion  2. Experimental procedure  3.1. Scale thickness and morphology  High purity (>99%) ZrB2 powder (\\x182 lm, Grade B, H.C. Starck, Newton, MA) was used to prepare the specimens for this study. Also, 2 wt% B4C (\\x180.8 lm, Grade HS, H.C. Starck) was added to all batches to remove the oxide impurities from the surface of ZrB2 powder particles and enhance densification [39]. For some batches, 4 mol% tungsten, molybdenum, or niobium was added in the form of W (\\x183 lm, Alfa Aesar, Ward Hill, MA), Mo (\\x183 lm, Alfa Aesar), or Nb (\\x181 lm, Alfa Aesar) powders all with reported purities of >99.9%. Hereafter, ZrB2 ceramics with only B4C additions are referred to as ‘‘nominally pure ZrB2’’, while ZrB2 with B4C and W, Mo, or Nb additions are referred to as (Zr,W)B2, (Zr,Mo)B2, and (Zr,Nb)B2, respectively. Ball milling in methyl ethyl ketone with ZrO2 media for 24 h was used to mix the powders. Measurement of the mass of zirconia media before and after ball milling showed that the amount of ZrO2 contamination added to the powder mixtures was less than 1 wt% based on the mass of ZrB2 powder. After drying and sieving (\\x0080 mesh), the powders were densified by hot pressing (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 2100 °C for 45 min at a pressure of 32 MPa. Bars with dimensions of 10 mm by 4 mm by 4 mm were cut from hot pressed billets that were 10 cm by 10 cm by 0.6 cm. The cut bars were finished by polishing sides with a final polishing step using a 15 lm diamond on all slurry. The bulk density of the hot pressed billets was measured by using the Archimedes technique, with water as the immersing medium following ASTM C373-88. X-ray diffraction (XRD; Philips X-Pert Pro diffractometer, Westborough, MA) was used to identify major crystalline phases present in the pre-oxidized composites. The measurements used Cu K-alpha radiation with a wavelength of 1.54 Å. The accelerating voltage was 45 kV, the tube current was 40 mA, the scan range was between 6° and 90° 2 theta, and the step size was 0.026°. A MoSi2 resistance-heated horizontal tube furnace (Model 0000543, CM Inc., Bloomfield, NJ), equipped with a high-purity alumina tube having a diameter of 6.35 cm, was used for oxidation studies. Specimens were heated to desired temperatures with a heating rate of \\x185 °C/min and an air flow rate of 0.2 cm/s (linear flow rate was calculated from the volumetric flow rate and the diameter of the tube). Specimens were removed from the furnace  All of the ZrB2 specimens reached relative densities of >95%. Archimedes’ density measurements indicated that open porosity was not significant in any of the specimens. Hence, porosity was not considered to have an adverse effect on the oxidation behavior. In addition, XRD analysis indicated that all ceramics contained a single phase that was either ZrB2 or a (Zr,TM)B2 solid solution, which is consistent with observations from previous studies [7,38]. Assuming that oxidation proceeds stoichiometrically, oxidation of (Zr,TM)B2 ceramics should produce a combination of molten B2O3 (melting temperature \\x18450 °C), solid ZrO2 (monoclinic below 1170 °C and tetragonal above), and solid TM-oxide. As reported in previous studies [7,38], EDS and XRD results showed that two distinct layers covered the surface of the ZrB2 and (Zr,TM)B2 ceramics after oxidation at temperatures between 800 and 1600 °C: (1) an outer glassy layer mainly consisting of B2O3, and (2) an inner layer mainly consisting of porous zirconia with the pores filled with B2O3. Due to low sensitivity of EDS to light elements, quantification of the boron content in the glassy phase was not possible by this method. However, results obtained by EDS indicated that the glassy phase contained O and TM along with a small amount of Zr. According to the ZrO2-B2O3 [40], WO3-B2O3 [41], and Nb2O5- B2O3 [42] phase diagrams, the solubility of ZrO2 and Nb2O5 in pure liquid B2O3 at 800 °C is negligible, while approximately 10 mol% WO3 can dissolve into B2O3 at that temperature. No phase diagrams were found for B2O3-Mo oxide systems. The formation of two-layer scales is due to immiscibility of ZrO2 and B2O3 combined with the large volume expansion (\\x18300% based on density calculations) associated with oxidation of ZrB2 to ZrO2 and B2O3 [6]. The addition of TMs to ZrB2 did not affect the formation of the two layer structure, but varying amounts of the TMs were incorporated into the B2O3, depending on the solubility of the particular TM oxide in B2O3 at the oxidation temperature. Fig. 1 shows SEM images obtained from the surfaces of nominally pure ZrB2, (Zr,W)B2, (Zr,Mo)B2, and (Zr,Nb)B2 after oxidation at 1600 °C for 0 h. The majority (80% or more of the area) of the surfaces of the specimens had light contrast, but some areas were darker. The light-colored areas contained ZrO2 and TM oxides, while the darker areas had a glassy appearance and was composed of B2O3 containing small amounts of Zr and TM. The amount of the  \\x0c', '226  M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  Fig. 1. Surfaces of (a) nominally pure ZrB2, (b) (Zr,W)B2, (c) (Zr,Mo)B2, and (d) (Zr,Nb)B2 after oxidation at 1600 °C for 0 h.  dark phase observed on the surfaces of ceramics oxidized at 1600 °C for 0 h increased from \\x185% surface coverage for nominally pure ZrB2 to \\x1810% for (Zr,W)B2, \\x1820% for (Zr,Mo)B2, and \\x1815% for (Zr,Nb)B2. At 1600 °C, the dissolution of ZrO2 into B2O3 increases to 24 mol% [40], while both WO3 and Nb2O5 are completely miscible with B2O3 due to the lower melting points of the TM oxides (1435 °C [41] for WO3 and 1485 °C [42] for Nb2O5) compared to the melting point of ZrO2 (2715 °C). Since MoO3 melts at 795 °C, it is also expected to be completely miscible with B2O3 at 1600 °C. The higher coverage of the glassy phase on (Zr,Nb)B2 compared to (Zr,W)B2 is consistent with evaporation data obtained by thermogravimetric analysis that showed lower weight loss from B2O3 containing Nb2O5 compared to B2O3 containing WO3 (data not included). One possible reason for the lower stability of the WO3-B2O3 glass is evaporation of WO3, which is volatile at temperatures above 1300 °C [43,44]. After oxidation at 1600 °C for 0 h, the thickness of the porous scale for nominally pure ZrB2 was \\x1885 lm compared to \\x1830 lm for (Zr,W)B2, and \\x1840 lm for both (Zr,Mo)B2 and (Zr,Nb)B2 (Table 1). The lower thickness of the scales on (Zr,TM)B2 ceramics is an indication of the effectiveness of W, Mo and Nb additions on improving the oxidation resistance of ZrB2. Since the liquid/glassy B2O3 scale is the diffusion barrier that protects unoxidized ZrB2 from oxygen in the external atmosphere, retention of the glassy phase on (Zr,TM)B2 specimens should improve the oxidation resistance of these ceramics compared to nominally pure ZrB2. Interestingly, (Zr,W)B2 had a lower porous scale thickness compared to (Zr,Mo)B2 and (Zr,Nb)B2, despite higher stability of the glassy phases on ZrB2 with Mo or Nb additions. The improved oxidation protection of the W-containing ceramic could be due to the effect of WO3, which alters the microstructure of the porous ZrO2 layer by liquid phase sintering [13,37]. The higher relative density of the ZrO2 scale on (Zr,W)B2 ceramics apparently resulted in lower  Table 1 Thickness of porous oxide scales of nominally pure ZrB2 and ZrB2 with 4 mol% W, Mo, or Nb after oxidation at 1600 °C for 0 (i.e., quenched as soon as it reached 1600 °C) or 3 h. The standard deviations represent at least 50 measurements performed on one cross section.  Material  ZrB2 (Zr,W)B2 (Zr,Mo)B2 (Zr,Nb)B2  Oxide scale thickness (lm)  0 h  85 ± 6 30 ± 3 40 ± 2 41 ± 2  3 h  570 ± 19 376 ± 16 386 ± 17 294 ± 13  therefore,  lower oxide scale thickness  oxygen permeability and, at this condition. After oxidation at 1600 °C for 3 h, the outer scales formed on all of the specimens consisted of a porous oxide layer. Some residual glassy phase was observed between the oxide particles in the (Zr,TM)B2 ceramics, but the glassy phase was not continuous. Although no continuous glassy layer covers the top surface, the glassy phase between the oxide particles could reduce permeability of the pore channels and could reduce oxidation rate [45]. The thickness of the porous oxide layer was highest for nominally pure (\\x18570 lm), while it was about 380 lm thick for ZrB2 (Zr,W)B2, about 390 lm thick for (Zr,Mo)B2, and about 300 lm thick for (Zr,Nb)B2 (Table 1). Hence, even after most of the glassy phase had evaporated, the transition metals had some beneficial effect on scale thickness. Fig. 2 shows the surfaces of nominally pure ZrB2 and (Zr,TM)B2 specimens after oxidation at 1600 °C for 3 h. The surfaces of nominally pure ZrB2 and (Zr,W)B2 were rough. As shown in the image insets in Fig. 2a and b, raised areas about 10 lm in diameter were observed on the surfaces. Charging around the highest points of the  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  227  Fig. 2. Surfaces of (a) nominally pure ZrB2, (b) (Zr,W)B2, (c) (Zr,Mo)B2, and (d) (Zr,Nb)B2 after oxidation at 1600 °C for 3 h. The black dashed lines shows the large raised areas on the surfaces, while the white dashed lines show the last areas of the surfaces that were covered by the glassy pools.  raised areas produced the contrast observed in the images. Some gaps were also observed between what appeared to be grains in the oxidized scale. The gaps were more severe at triple points between the oxide grains on the surfaces of ZrB2 and (Zr,W)B2. The similarity of the morphology of the oxide scales for nominally pure ZrB2 and (Zr,W)B2 at 1600 °C may be due to the volatility of the respective liquid oxides. Both B2O3 and WO3 are volatile at 1600 °C. As concluded from a previous study on the oxidation of (Zr,Nb)B2 ceramics at lower temperatures [21], the complete evaporation of these oxides from the surface, along with the accompanying precipitation of ZrO2 from the evaporating liquid, may be a contributing factor to the morphology of the surface of the scale. The surfaces of (Zr,Mo)B2 and (Zr,Nb)B2 were different than nominally pure ZrB2 and (Zr,W)B2 after oxidation at 1600 °C for 3 h. The white dashed lines in Fig. 2c and d correspond to the last areas of the surfaces that were covered by pools of liquid oxide. Bubbles and ruptured bubbles were observed in some of the areas between the glassy pools as shown by black dashed lines. The effect was more pronounced for (Zr,Mo)B2 than (Zr,Nb)B2. As shown in the image insets in Fig. 2c and d, the areas between the pools on (Zr,Mo)B2 and (Zr,Nb)B2 were continuous and showed no sign of rupture or gaps. The areas between the pools appeared to be nearly fully dense with the porosity of the surrounding oxide layer increasing near the pools. Hence, the lower thickness of the  porous scales on (Zr,Mo)B2 and (Zr,Nb)B2 may be due to the higher apparent density of the external scale and/or the retention of more glassy oxide during oxidation. Fig. 3 shows cross sectional SEM images from selected areas in scales that formed on nominally pure ZrB2 and (Zr,Nb)B2 after oxidation at 1600 °C for 3 h. The feature marked with an arrow in Fig. 3a is similar to the high points of raised areas in Fig. 2a. The scale appears to be thicker in this area. In addition, the gap indicates poor adhesion between the outer layer of the scale and the layer beneath it. The raised areas were thought to form due to evaporation of B2O3 from within the scale. Fig. 3b shows a cross section of an area near what appears to be a void in the oxide scale. This area also seems to be where the scale has delaminated. Similar voids are seen in the oxide scale on both (Zr,Mo)B2 and (Zr,Nb)B2. Presumably, the voids are channels to allow release of B2O3 vapor, which cannot otherwise escape as the surrounding scale appears to be dense. The voids may be caused by rupture of the top surface of the oxide layers on (Zr,Mo)B2 and (Zr,Nb)B2. These ruptures appear to be similar to those observed by Karlsdottir et al. [40], which occurred as result of evaporation of SiO(g) formed during oxidation of ZrB2-SiC ceramics. However, since the scales on nominally pure ZrB2 and (Zr,W)B2 are composed of crystalline ZrO2-based grains, rupture appears to occur at grain boundaries, resulting in several small  Fig. 3. Fracture surfaces of (a) nominally pure ZrB2, and (b) (Zr,Nb)B2 after oxidation at 1600 °C for 3 h, showing the difference between the raised areas on ZrB2 and (Zr,Nb)B2.  \\x0c', '228  M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  ruptures across the surfaces. As a result, the outer oxide layers on nominally pure ZrB2 and (Zr,W)B2 are not adhered well to the layers beneath and tend to spall. The top layers on (Zr,Mo)B2 and (Zr,Nb)B2 have better adhesion and less tendency to spall.  3.2. Dark and light zirconia scale layers  Previously published studies, and the analysis presented above, describe oxide scale thickness as a quantity that can be readily measured. However, examination as part of the present study has revealed several complications that have likely led to inconsistencies when measuring and reporting scale thicknesses. The next several paragraphs describe the presence of two sub-layers within what is described as the ‘‘porous ZrO2’’ layer above. Specimen size raises additional complications with measurements that are discussed. Fig. 4 shows optical micrographs of fracture surfaces of nominally pure ZrB2 and (Zr,TM)B2 specimens after oxidation at 1600 °C for 6 h. Here, specimens were oxidized at 1600 °C for 6 h to thicken both layers and enable better resolution using optical microscopy. The oxide scales consisted of two layers similar to observations of Zhang et al. [37]. The outer layer (outside the white line) was lighter in contrast while the inner layer was darker. The dark layer appeared dense with a metallic shine, like the unoxidized matrix. In contrast, the outer layer had a lighter color, was porous, and did not exhibit a metallic appearance. The dark layer appeared to be strongly adhered to the unoxidized matrix with no delamination observed between the dark layer and the matrix. However, adhesion between the light and dark layers was weaker, which often resulted in damage to the interface during polishing. Studies of the oxidation of zirconium have reported the presences of a dark oxide layer that is dense, adherent, and protective, resulting in parabolic oxidation kinetics [46], which indicate that the oxide layer acts as a barrier to inward transport (presumably diffusion) of oxygen. For Zr and its alloys, a process called ‘‘transition’’ or ‘‘breakaway’’ converts part of the dark layer (presumably oxygen-deficient ZrO2) to a lighter layer that is porous and tends to spall. Once this transformation occurs, linear oxidation kinetics are observed, indicating that the oxide layer is no longer protective. The mechanism of breakaway is still controversial [46,47]. The formation of a similar dark scale has been reported for ZrB2 ceramics [20,37,48], but has not been discussed in detail in any publication. Higher apparent density of the dark layer compared to the light layer should result in lower oxygen transport through the dark  layer, which might make it the rate limiting step for oxidation in the absence of the liquid/glassy B2O3 layer. To highlight the difficulties associated with identifying and measuring the oxide scale thickness, Fig. 5 shows SEM images of a fracture surface and polished cross section of (Zr,W)B2 after oxidation at 1600 °C for 6 h. The dark layer is nearly impossible to distinguish from the unoxidized matrix in SEM images of polished specimens. Hence, the thickness of the dark layer is likely to not be included in scale thickness measurements for studies that utilize polished cross sections and SEM images for analysis of oxidation behavior. Examination of fracture surfaces results in more contrast due to topology differences between the dark layer and the unoxidized matrix that are removed by polishing; however, the dark and light layers were not distinguishable from the fracture surface. As a result, the scale thickness might be reported as \\x181000 lm from the fracture surface, but \\x18540 lm from the polished section even though these specimens were oxidized under the same conditions. This discrepancy is only an issue for oxidation conditions at which the glassy scale has retreated beneath the outer surface. The best ways to characterize scale thicknesses, to capture both the dark and light layers, seems to be optical microscopy combined with SEM of fractured surfaces. Regions of the oxide scale from which the glassy phase has evaporated appear lighter in color (i.e., white and porous) while regions in which the pores are still filled with B2O3 appear dark. Studies using SEM analysis of polished sections may not observe the layer that appears darker by optical microscopy and, therefore, underestimate the total oxide layer thickness. Another complication is that oxide layers can have a curved shape after severe oxidation as shown in Figs. 4 and 5. The curvature is an artifact of specimen size. In this case, the oxide-matrix interface curves (i.e., the oxide layer becomes thicker) near the corners of the original specimen due to oxygen permeation from orthogonal faces. For the present study, all the measurements of oxide scale thicknesses were performed near the middle of the faces of specimens, which had lower curvatures and were assumed to be more representative of the oxidation of an infinite plate. Curvature becomes a significant problem when the thickness of the oxide layers is of the same order of magnitude at the specimen dimensions (i.e., oxide thickness is \\x18500 lm for a 5 mm wide specimen. Considering oxidation at  thickness after for (Zr,W)B2 at  scale, the highest  the was  both 1600 °C  layers for  of 6 h  Fig. 4. Optical micrographs of fracture surfaces of (a) nominally pure ZrB2, (b) (Zr,W)B2, (c) (Zr,Mo)B2, and (d) (Zr,Nb)B2 after oxidation at 1600 °C for 6 h.  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  229  Fig. 5. SEM images of (a) a fracture surface, and (b) a polished cross section of (Zr,W)B2 after oxidation at 1600 °C for 6 h.  910 ± 37 lm followed by 860 ± 34 lm for ZrB2, 610 ± 27 lm for (Zr,Nb)B2, and 520 ± 22 lm for (Zr,Mo)B2. The higher apparent density of the dark scale should make it a better barrier to oxidation compared to the lighter scale. Therefore, one method to judge the vulnerability of specimens to oxidation is the thickness of the lighter scale layer. Nominally pure ZrB2 had the greatest thickness of the light layer (700 ± 25 lm), followed by (Zr,W)B2 (460 ± 18 lm). The lighter layers were significantly thinner for (Zr,Nb)B2 (180 ± 10 lm) and (Zr,Mo)B2 (120 ± 7 lm). Likewise, the dark layer was thinnest for nominally pure ZrB2 (160 ± 9 lm) compared to dark layers that were more than twice as thick for (Zr,Mo)B2 (400 ± 15 lm), (Zr,Nb)B2 (430 ± 17 lm), and (Zr,W)B2 (450 ± 19 lm). While W additions may not provide a significant improvement to oxidation resistance at 1600 °C, Mo and Nb additions resulted in significantly lower thicknesses for the light oxide layers. As mentioned above, the light layer forms as B2O3 evaporates. The lower thickness of the light layer for (Zr,Mo)B2, (Zr,Nb)B2, and (Zr,W)B2 compared to nominally pure ZrB2 indicates a higher stability for the glassy phases in these materials, resulting in thicker dark oxide layers. Hence, (Zr,Mo)B2 and (Zr,Nb)B2 have the best oxidation resistance based on lowest overall scale thickness. Further, the dark layers in these compositions appear to be more effective barriers to oxygen transport, which results in lower light scale layer thicknesses compared to (Zr,W)B2. As discussed in a previous article [38], the outer surface of the porous oxide scale is initially composed of equiaxed particles when it is still covered by a layer of liquid/glassy B2O3. The particles form as the result of dissolution of ZrO2 and TM oxides into the B2O3 melt followed by reprecipitation. As the ZrO2 particles begin to form, they are surrounded by liquid B2O3 and grow uniformly. As B2O3 evaporates, the solution-precipitation mechanism changes due to loss of liquid B2O3 at the outer surface. Although ZrO2 continues to precipitate as the B2O3 scale evaporates, it can only grow  where the ZrO2 particles are in contact with the liquid phase. As the B2O3 surface recedes into the porous oxide, ZrO2 takes a columnar morphology, presumably due to preferential growth from the liquid interface as it recedes into the oxide scale. The black dashed lines in Fig. 4 indicate the border between equiaxed ZrO2 and ZrO2 with the columnar morphology for the scale on nominally pure ZrB2 and (Zr,W)B2. Under the same oxidation conditions, the light layer on (Zr,Mo)B2 and (Zr,Nb)B2 specimens was composed of only equiaxed particles. Liquid phases with higher stability at lower oxidation temperatures resulted in thicker primary precipitated oxide layers on (Zr,Mo)B2 and (Zr,Nb)B2. In turn, the thicker primary oxide layers may lead to improved oxidation resistance due to the more uniform morphology. In addition, the equiaxed oxide that is more prevalent on (Zr,Mo)B2 and (Zr,Nb)B2 may be less prone to spallation than the columnar oxide on ZrB2 based on observations described above.  3.3. Model of oxidation of ZrB2 ceramics at 1600 °C  Fig. 6 is a schematic that describes evolution of the structure of the oxide scale that forms during oxidation of ZrB2 ceramics at 1600 °C. In the early stages of oxidation (left panel), a liquid B2O3 layer is present along with ZrO2 particles. In this stage, a continuous liquid B2O3 layer is present and liquid B2O3 fills all of the volume between ZrO2 particles. In optical microscopy, this oxide scale has an iridescent appearance, similar to observations of scales produced by oxidation of zirconium and zirconium alloys [46,49,50]. In the initial stages of oxidation, equiaxed oxide particles form beneath the liquid B2O3 layer. The particles form by solution-precipitation as has been discussed elsewhere [38]. With increasing oxidation temperature or time, evaporation of B2O3 increases and some parts of the porous layer are exposed, and, consequently, become the layer that appears lighter in optical micrographs.  Fig. 6. Model of the evolution of the oxide structure of ZrB2 ceramics during oxidation at 1600 °C.  \\x0c', '230  M. Kazemzadeh Dehdashti et al. / Corrosion Science 91 (2015) 224-231  Acknowledgments  This work was supported as part of the National Hypersonic Science Center for Materials and Structures (Grant FA9550-09-10477) with Dr. Ali Sayir (AFOSR) and Dr. Anthony Calomino (NASA) as program managers. The authors wish to thank project principal investigator Dr. David Marshall of Teledyne Scientific and Imaging for his support and guidance.  References  Evaporation of B2O3 is more severe at higher temperatures, resulting in recession of the liquid into the porous layer (second panel in the schematic), and, eventually, nearly complete loss of B2O3 (final panel in the schematic). When the liquid recedes from the outer surface, ZrO2 particles can no longer form at the outer surface or grow uniformly due to a lack of liquid phase to transport dissolved ZrO2. Continued growth of ZrO2 presumably occurs mainly from the receding front, which is now beneath the outer surface of the specimen. Hence, the columnar morphology may develop as the originally equiaxed particles act as growth sites for ZrO2 precipitating from the receding B2O3. At temperatures at which the (TM-oxide)-B2O3 layer formed on (Zr,TM)B2 also evaporates (i.e. above about 1400 °C) the differences between the different TM additions were more obvious. The higher volatility of WO3 and the resulting WO3-B2O3 mixtures, compared to Nb2O5-B2O3, resulted in lower coverage of the surface of (Zr,W)B2 by the glassy phase compared to (Zr,Nb)B2 after oxidation at 1600 °C for 0 h. The thicker porous zirconia layer for (Zr,Nb)B2 compared to (Zr,W)B2 was presumably due to liquid phase sintering of ZrO2 in the presence of WO3. However, increasing the oxidation time to 3 h at 1600 °C resulted in more WO3 evaporation making the surface of (Zr,W)B2 appear similar to nominally pure ZrB2 rather than (Zr,Nb)B2. Earlier evaporation of the protective glassy layer from (Zr,W)B2 resulted in a thicker porous oxide layer for (Zr,W)B2 compared to (Zr,Nb)B2. Although no phase diagrams were available for the MoO3-B2O3 and ZrO2-MoO3 systems, the oxidation behavior of (Zr,Mo)B2 was similar to (Zr,Nb)B2, meaning that the MoO3-B2O3 and ZrO2-MoO3 systems resemble their Nb analogues. After oxidation at 1600 °C for 6 h, the porous (light) oxide layer was thicker for (Zr,W)B2 compared to (Zr,Mo)B2 or (Zr,Nb)B2. However, the thicknesses of the dark oxide layers were almost the same for all of the (Zr,TM)B2 ceramics and higher than for nominally pure ZrB2. It can be speculated that all of the TM additions have about the same effect on the transport rate of oxygen from the interface between the light and dark layers to the unoxidized matrix. The larger thickness of the light scale layer in (Zr,W)B2 is due to the higher volatility of WO3, which resulted in earlier loss of the protective liquid layer and more oxidation compared to (Zr,Mo)B2 and (Zr,Nb)B2.  4. Conclusion  The effects of additions of W, Mo, or Nb on the oxidation behavior of ZrB2 were investigated. After oxidation at 1600 °C for 0 h, the higher stability of Nb2O5-B2O3 and MoO3-B2O3 resulted in lower oxide scale thicknesses for (Zr,Nb)B2 and (Zr,Mo)B2 compared to nominally pure ZrB2 and (Zr,W)B2 specimens. Two distinct layers were observed in the oxide scale after oxidation at 1600 °C for 6 h, an outer layer that was lighter in color and an inner layer that was darker. The thickness of the light layer for nominally pure ZrB2 was almost 1.5 times greater than for (Zr,W)B2 and 4 times higher than (Zr,Mo)B2 and (Zr,Nb)B2. Lower scale thicknesses for the light layer for (Zr,Mo)B2 and (Zr,Nb)B2 indicated that Mo and Nb were more effective in reducing oxygen transport from the external atmosphere toward the interface between the light and dark layers. However, the thicknesses of the dark layers were about the same for all of the (Zr,TM)B2 specimens and almost twice as thick as the dark layer for nominally pure ZrB2, showing the same effect of TM additions on lowering oxygen diffusion through the dark layer toward the unoxidized matrix. Additions of Mo or Nb were most effective in improving the oxidation resistance of ZrB2 by increasing the stability of the protective liquid/glassy B2O3 layer and reducing oxygen transport through the light layer.  [20]  air,  J.  [11]  in dry  [1] K. Upadhya, J.M. Yang, W.P. 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},{
  "_id": 66,
  "PDF": "Effects of ZrB2 and SiC dual addition on the oxidation resistance of graphite at 1600–2000°C.pdf",
  "Text": "['Corrosion Science 76 (2013) 182-191  Contents lists available at SciVerse ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Effects of ZrB2 and SiC dual addition on the oxidation resistance of graphite at 1600-2000 °C  Zeng-Hua Gao a,  Jing-Jun Xu a, Zhong-Wei Zhang b, Yu-Hai Qian a, Mei-Shuan Li a,⇑  a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China b Aerospace Research Institute of Materials and Processing Technology, Beijing 100076, China  a r t i c l e  i n f o  a b s t r a c t  The effects of ZrB2 and ZrB2 + SiC additions on the oxidation kinetics of graphite at 1600-2000 °C in air were investigated. The ZrB2 + SiC dual addition improves the oxidation resistance of graphite more effectively than the ZrB2 single addition, because the oxide scale formed on C-ZrB2-SiC is denser and thinner due to the existence of glassy SiO2. As the oxidation temperature increases, the oxidation rate of C-ZrB2- SiC gradually increases and oxide scales with layered microstructures form on its surface due to the greatly enhanced active oxidation of SiC at higher temperatures. Ó 2013 Elsevier Ltd. All rights reserved.  Article history:  Received 10 May 2013 Accepted 24 June 2013 Available online 2 July 2013  Keywords:  B. Weight loss B. XRD B. SEM C. Oxidation C. Kinetic parameters C. Thermodynamic diagrams  1. Introduction  Exploration and selection of novel ultra-high temperature materials (UHTMs) are tremendously driven by the demand of advanced thermal protection systems (TPS) for use in the extreme environments (above 1400 °C in oxidizing atmospheres) related to hypersonic flight, atmospheric re-entry and rocket propulsion [1-3]. Graphite and graphite-based composite materials have drawn great attention due to their exceptional characteristics including low density, high melting temperature, high specific strength and modulus, high thermal conductivity, excellent thermal shock resistance, and good machinability [4,5]. However, graphite easily burns itself away by reacting rapidly with oxygen at temperatures above 500 °C, which causes serious degradation of the thermal and mechanical properties of the graphite materials and hence limits their applications at high temperatures [6-8]. Numerous studies have been conducted in an attempt to protect graphite or graphite-based composites against oxidation at high temperatures. The introducing of Zr-containing ultra-high temperature ceramics (UHTCs) into graphite or C/C composite, such as ZrC [9-11], ZrC-TaC [12], ZrB2 [13], ZrB2-B4C [14], was reported to be an effective approach due to the formation of a ZrO2 scale during oxidation. It was generally concluded in those work that, the ZrO2  ⇑ Corresponding author. Address: 72 Wenhua Road, Shenhe District, Shenyang 110016, China. Tel.: +86 24 23971145; fax: +86 24 23891320.  E-mail address: mshli@imr.ac.cn (M.-S. Li).  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.06.041  scale could act as a barrier for both heat transfer and oxygen transport to the underneath substrate because of its low thermal conductivity and oxygen permeability. Nonetheless, the ZrO2 scale formed during the oxidation of graphite with single Zr-containing addition tends to be porous and discontinuous, especially when the addition content is low. Actually, even for the monolithic ZrB2 ceramic, its active oxidation occurs above 1400 °C due to the intensive evaporation of B2O3, which leaves behind a porous, non-protective ZrO2 layer [15,16]. However, it has been well known that the addition of SiC could improve significantly the oxidation resistance of ZrB2 at higher temperatures by encouraging the formation of borosilicate glass on the exposed surface, which has a higher boiling point, higher viscosity and lower vapour pressure than B2O3 [17-25]. Therefore, it is expected that the high temperature oxidation resistance of graphite can be further improved by adding both ZrB2 and SiC as oxidation inhibitors. Very recently, Liu and Li et al. investigated the effects of ZrB2-ZrC-SiC [26] and ZrB2-SiC [27] additions on the ablation resistance of C/C composite using oxyacetylene torch with different heat fluxes. It was concluded that the extra SiC addition could enhance the ablation resistance of C/ C composite up to 2300 °C. But the mechanisms of the effects of ZrB2 and SiC on the oxidation resistance of carbon materials are still not well understood. On the other hand, oxyacetylene torches and arc heaters are two kinds of commonly used techniques for evaluating the ablation resistance of materials above 1600 °C. Though it simulates well the practical conditions encountered in aerospace applications, the environments produced by these two testing techniques  \\x0c', 'Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  183  are not well controllable (such as temperature, oxidizing species composition ) and cause the removal of material by a combination of thermo-physical, thermo-chemical, and thermo-mechanical factors [14,27,28]. In this way, the degradation mechanisms of materials during testing are interacted and not well distinguishable. To identify specifically the isothermal oxidation behaviour of materials at ultra-high temperatures (>1600 °C), new testing techniques, which are capable of producing well controllable ultra-high temperature oxidizing environments, are desired. In present work, first, we introduced the induction heatingbased ultra-high temperature oxidation testing apparatus developed in our laboratory. Using this apparatus, the effects of ZrB2 and ZrB2 + SiC additions on the isothermal oxidation behaviour of graphite at 1600-2000 °C up to 1800 s in air were investigated and compared. The oxidation kinetics were characterized by the in situ monitored apparent radius loss of the cylindrical sample during oxidation as well as the ex situ measured weight loss after oxidation.  2. Experimental  2.1. Materials and preparation  Pure graphite (denoted as G), and two graphite-based composites, C-ZrB2 and C-ZrB2-SiC (denoted as GZ and GZS, respectively), were supplied by Key Laboratory of Carbon Materials, Institute of Coal Chemistry, Chinese Academy of Sciences. The composite materials were synthesized by hot-pressing. The starting materials, including commercial ZrB2, SiC and graphite powders, are 99% in purity and \\x1810 m in size. Graphite powders were mixed with different ceramic powders in certain ratios, as listed in Table 1. The powders mixture was milled for 24 h in alcohol by attrition milling with agate media. Then, the slurry was dried at 60 °C for 12 h in an oven. The mixed powders were hot-pressed under a pressure of 50 MPa in a graphite mould at 2000 °C for 1 h in flowing Ar. At last, the composites were graphitized at 2500 °C. Some properties of  Table 1 Nominal compositions and some properties of three materials.  Designation  Composition (wt.%)  Bulk density (g/cm3)  Porosity (%)  G GZ GZS  100C 82C + 18ZrB2 76C + 18ZrB2 + 6SiC  1.72 2.18 2.28  12 12 11  three materials are also listed in Table 1. Cylindrical samples (20 mm in diameter, 20 mm in height) for the oxidation tests were cut from the as-prepared bulk materials.  2.2. Ultra-high temperature oxidation testing apparatus  To specifically study the isothermal oxidation behaviour of materials at temperatures above 1600 °C, an ultra-high temperature oxidation testing apparatus based on the induction heating technology was developed in our laboratory, as schematically shown in Fig. 1a. It is originally assembled by an intermediate frequency power supply (20 kHz/30 kW) with a water-cooled copper induction coil, a testing chamber and vacuum pumps of a mechanical pump and a Roots pump. The surface temperature of the sample is measured by a two colour ratio pyrometer (Heitronics KT 18.03 R, 1000-3000 °C) and can be well controlled by adjusting the output power of the power supply. It should be noted that induction coils with various shapes and dimensions are available for samples with different configurations. Details of the physical principles and the temperature calibration of this induction heating facility can be found elsewhere [12,29]. The ambient pressure and flow rate in the testing chamber can be adjusted and maintained constant by controlling the exit pumping speed and the volume of gas inlet through the leak valve. What’s more, the apparatus has recently been modified by equipping with a micro-vision monitoring camera, from which the dimensional changes of the sample during oxidation testing can be in situ observed and photographically recorded at a resolution of ± 0.02 mm and at frequent intervals.  2.3. Oxidation testing procedure  In the oxidation test, the cylindrical sample was placed on a graphite susceptor positioned in the centre of the induction coil. The testing chamber was evacuated to a pressure of \\x182 Pa (i.e. \\x182 \\x02 10\\x005 atm) prior to heating. Then, the sample was heated at an average heating speed of \\x1820 °C/s to the target temperature, and a 5 min isothermal hold is followed. Afterwards, air was rapidly introduced into the testing chamber. Normally, the ambient air pressure can be achieved in less than 10 s. Meanwhile, the in situ monitoring of the sample by the micro-vision camera was activated. The surface temperature of the sample was maintained constant at the target temperature during the testing period. At the end of the oxidation test, the power supply of the inductor was turned off and the testing chamber was evacuated again to a  Fig. 1. (a) Illustration of the induction heating-based ultra-high temperature oxidation testing apparatus, (b) a self-heated sample, and (c) sample morphologies monitored by the micro-vision camera during oxidation at different time.  \\x0c', '184  Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  pressure of \\x182 Pa. The sample was first cooled down to around 1000 °C at a rate of \\x1810 °C/s, and then to room temperature in \\x1815 min. Fig. 1b shows the appearance of a self-heated sample placed in the centre of a water-cooled copper induction coil.  2.4. Oxidation kinetics characterization  In the oxidation test, the in situ appearance of the sample was continuously photographed by the micro-vision camera at a high magnification and at intervals of 1 s. Fig. 1c shows a set of pictures of the incandescent sample monitored in situ at different time during oxidation testing. From which, the apparent radial dimension changes of the sample during oxidation were determined and plotted as curves of the apparent radius loss as a function of the oxidation time. This characterization method has long been applied in the studies on the oxidation and combustion of carbon [30,31]. The dash line in Fig. 1c indicates the location from where the radial dimension of the sample was determined. Apart from the in situ monitored apparent radius loss of cylindrical sample, the weight loss per unit surface area of sample after oxidation, Dw/A, was also determined ex comparison.  situ for  the the  Dw  A  ¼ w0 \\x00 wt A  ð1Þ  where w0 and wt are the weight of the cylindrical sample before and after oxidation, respectively; is the oxidation time; A is the exposed surface area of the cylindrical sample. The weight measurements were conducted using an electronic balance with an accuracy of 10\\x005 g.  t  2.5. Surface characterization  Phase compositions of the surface scales of the as-oxidized composite samples were characterized by an X-ray diffractometer (XRD, Rigaku D/max-2400, Tokyo, Japan) with Cu Ka radiation. The surface scales were mechanically peeled off from the sample surface and ground to powders in an agate mortar before XRD analysis. Surface and cross section morphologies of the as-oxidized composite samples were investigated using a SUPRA 35 scanning electron microscope (SEM, LEO, Oberkochen, Germany) equipped with energy-dispersive spectroscopy (EDS, Oxford Instrument, Oxford, U.K.). To observe the cross section, the as-oxidized samples were first mounted in resin and then cut using a diamond wafering blade. The cut-out cross sections were ground down to 2000-grit SiC paper and polished with 1.5-lm diamond paste. Prior to SEM observation, a thin layer of Au was sputtered on the surface of the as-oxidized samples to avoid electron charging. The plane porosity of the oxide scales and the pore sizes were determined from SEM images using image analysis software (Image J, National Institutes of Health, Bethesda, MD).  3. Results  3.1. Oxidation kinetics  Plots of the apparent radius loss as a function of the oxidation time for three materials exposed at 1600 °C in air up to 1800 s are shown in Fig. 2a. The apparent radius losses of G, GZ and GZS all follow the linear rate law at this temperature. However, GZ and GZS have much lower apparent radius loss rates, especially in the case of ZrB2 + SiC dual addition. It indicates that, the dual addition of ZrB2 + SiC enhances the dimensional integrity of graphite in ultra-high temperature oxidizing atmosphere more effectively than the single addition of ZrB2.  Fig. 2. In situ monitored apparent radius loss during the oxidation of (a) G, GZ and GZS at 1600 °C, (b) GZS at 1600-2000 °C.  Fig. 2b shows the curves of the apparent radius loss as a function of the oxidation time for GZS exposed at 1800 and 2000 °C. Generally, they display similar upward trends with that at 1600 °C. But unlike the distinct linear characteristic all along for the oxidation at 1600 °C, the apparent radius loss curves for the samples exposed at 1800 and 2000 °C appear several fluctuations. At an early stage of 0-300 s, the sample exposed at 1800 °C exhibits comparable radius losses with that exposed at 1600 °C. However, the apparent radius loss for the oxidation at 1800 °C experiences several sharp inflections between 300 and 960 s and then rises linearly at a higher rate than that at 1600 °C. While for the oxidation at 2000 °C, the apparent radius loss rises rapidly from the very beginning of the oxidation. Besides, it experiences several mild inflections before 1360 s and then rises linearly at a similar rate with that at 1800 °C. The fluctuations of the apparent radius loss suggest that the inflation and/or disruption of the surface scales on the samples occurred during oxidation at 1800 and 2000 °C. For comparison, the weight loss results of three materials exposed at 1600 °C and GZS exposed at 1600-2000 °C are plotted in Fig. 3a and b, respectively. It can be seen that the weight losses of three materials at 1600 °C also increase linearly with increasing the oxidation time. However, the composite materials display much slower weight loss rates than pure graphite, and GZS exhibits the lowest weight loss rate. These results are in well accordance with the in situ monitored apparent radius loss characteristics shown in Fig. 2a. While for the oxidation of GZS at 1800 and 2000 °C, the weight losses also display linear characteristics but develop at higher rates, as shown in Fig. 3b. It should also be noted that, the weight losses of the samples after oxidation at 1600 and 1800 °C for 300 s are quite close, which also agrees well with the apparent radius loss characteristics shown in Fig. 2b.  \\x0c', 'Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  185  Fig. 3. Weight loss during the oxidation of (a) G, GZ and GZS at 1600 °C, (b) GZS at 1600-2000 °C. Data points are experimental results and solid lines are fitting results.  Fig. 4. XRD patterns of (a) the surface scales formed on GZ and GZS after oxidation at 1600-2000 °C for 1800 s and (b) the powdery deposits on the wall of the testing chamber after the oxidation testing of GZS at 2000 °C.  Table 2 Oxidation rates of three materials at 1600-2000 °C.  Material  T (°C)  Apparent radius loss rate  Weight loss rate  G  GZ GZS  Kr (mm/s)  0.9 \\x02 10\\x003  1.6 \\x02 10\\x004 1.0 \\x02 10\\x004 - -  1600  1600 1600 1800 2000  R2 (%)  Kw (g/cm2\\x01s)  R2 (%)  99.6  99.7 98.2 - -  1.7 \\x02 10\\x004 1.5 \\x02 10\\x004a 7.6 \\x02 10\\x005 4.1 \\x02 10\\x005 8.9 \\x02 10\\x005 1.1 \\x02 10\\x004  99.2 99.9 98.8 99.8 99.5 96.7  a Derived from the corresponding apparent radius loss data in Fig. 2a.  By fitting the experimental results in Fig. 2a and Fig. 3, the corresponding apparent radius loss rate (Kr) and weight loss rate (Kw) of three materials at different temperatures were obtained, as listed in Table 2. It can be seen that the apparent radius loss rate of GZS during oxidation at 1600 °C, which is 1.0 \\x02 10\\x004 mm/s, is almost an order of magnitude lower than that of pure graphite and that of GZ, which are 1.6 \\x02 10\\x004 and 9.0 \\x02 10\\x004 mm/s, 63% of respectively. On the other hand, the weight loss rate of GZS during oxidation at 1600 °C, which is 4.1 \\x02 10\\x005 g/cm2 s, is about a quarter of that of pure graphite and half of that of GZ, which are 1.7 \\x02 10\\x004 and 7.6 \\x02 10\\x005 g/cm2 s, respectively. For the oxidation of GZS at 1800 and 2000 °C, it is not reliable to obtain the radius loss rates by simply fitting the in situ monitored apparent radius loss curves linearly due to their fluctuant characteristics. But those curves supply information about the inflation and disruption of the surface scales on the samples during oxidation at those temperatures. On the other hand, the weight loss rates of GZS at 1800 and 2000 °C were determined to be  8.9 \\x02 10\\x005 and 1.1 \\x02 10\\x004 g/cm2 s, respectively. It indicates that the weight loss rate of GZS increases gradually with increasing the oxidation temperature. However, it is still lower than that of graphite at 1600 °C, as listed in Table 2.  3.2.  Identification of the oxidation products  No condensed oxide phase existed on the surface of the as-oxidized pure graphite sample. While for the composite samples, oxide scales formed on their surface during oxidation. Fig. 4a shows the XRD patterns of the surface scales of the composite samples after oxidation at different temperatures for 1800 s. For the oxidation of GZ and GZS at 1600 °C, only m-ZrO2 was detected in the surface scale. While for GZS exposed at 1800 °C, m-ZrO2 still existed as the major phase. However, ZrB2 also appeared at 1800 °C and became the major phase in the surface scale at 2000 °C. This implies the existence of a ZrB2 transition layer, Cand SiC-depleted layer, between the oxide scale and the unaffected substrate. In all cases, no boronand/or silicon-containing crystalline oxide phase has been detected in the surface scales by XRD analysis. However, white powdery deposits were found on the wall of the testing chamber after the oxidation testing of GZS, and their amount increased with the increase of the testing temperature. The XRD pattern in Fig. 4b shows that the deposits are composed of partially crystallized B2O3 and SiO2, suggesting that evaporations of boronand silicon-containing oxides took place during the oxidation of GZS. White powdery deposits were also found after the oxidation testing of GZ at 1600 °C and identified to be partially crystallized B2O3 (the XRD result is not shown here for brevity).  i.e. a  \\x0c', '186  Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  The composite samples before oxidation have slightly different colours but the same size with the pure graphite sample. It can be seen that, after oxidation at 1600 °C, the graphite sample (labelled as G/1600 °C) exhibits an obvious decreased size and a rough surface, while the composite samples are covered entirely by white scales and show slightly changed sizes. Fig. 6 shows the surface and cross section morphologies of the oxide scales formed on GZ and GZS samples after oxidation at 1600 °C for 1800 s. The dark areas in all pictures represent the locations of pores. It should be noted that the pores in the surface scale were filled with resin during the mounting procedure for cross section observations (see Fig. 6a0 and b0 ). The elemental compositions of all the phases observed were identified by EDS analysis. For GZ sample (see Fig. 6a and a0 ), a single-layer oxide scale with highly porous and interconnected skeleton microstructure formed on its surface (the white phase). This is similar to the microstructure of the oxide scale formed on ZrB2/C composites reported by Li et al. [13]. According to the XRD result (Fig. 4a), the oxide scale is composed of ZrO2 grains. The EDS analysis confirmed the presence of Zr and O elements in the oxide scale. The ZrO2 grains have an average size of \\x188.6 lm [see the high magnification inset in Fig. 6b]. The pores distributing in and between the porous ZrO2 skeleton have sizes of \\x185 and \\x18200 lm, the average respectively. The plane porosity of the oxide scale was determined to be \\x1850%. The thickness of the oxide scale is \\x18700 lm. For GZS sample (see Fig. 6b and b0 ), the oxide scale formed on its surface also has a single-layer skeleton microstructure. However, it is denser due to the existence of a grey glassy phase [see the high magnification inset in Fig. 6b]. Since only Si and O elements were detected by EDS, the grey phase was believed to be amorphous SiO2. The ZrO2 grains (the white phase) are embedded in the amorphous SiO2 and have a smaller  Fig. 5. Macrographs of various samples, (a) graphite before oxidation, (b) graphite after oxidation at 1600 °C for 1500 s, (c-f) graphite-based composites after oxidation at 1600-2000 °C for 1800 s.  3.3. Morphological observations of the surface scales  Macrographs of a cylindrical pure graphite sample before and after oxidation at 1600 °C for 1500 s, and composite samples after oxidation at different temperatures for 1800 s are shown in Fig. 5.  Fig. 6. Surface and cross section morphologies of (a, a0 ) GZ and (b, b0 ) GZS after oxidation at 1600 °C for 1800 s.  \\x0c', 'Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  187  average size of \\x183.4 lm. The pores mainly locate between the dense ZrO2/SiO2 skeleton with an average size of \\x18200 lm. A similar porous oxide scale was observed by Tang et al. [32] on an asablated C/C-ZrB2-SiC composite. The plane porosity of the oxide scale was determined to be \\x1830%, which is 40% lower than that of GZ. Besides, the oxide scale has a much smaller thickness of \\x18225 lm, which is less than a third of that on GZ. It should be noted that, no boron has been detected in the oxide scales for both composite materials by EDS. Fig. 7a and b shows the surface morphologies of the oxide scales formed on GZS samples after oxidation at 1800 and 2000 °C for 1800 s. It can be seen that, more glassy SiO2 (the grey phase) was produced on the outer surface of the oxide scales. Bulges and pores were also observed and their amounts increased with the increase of temperature. The cross section morphologies of these scales were shown in Fig. 7a0 and b0 . It should be noted that the oxide scales formed at 1800 and 2000 °C adhered to the underneath substrates in origin, but they were separated by resin during the mounting procedure for the cross section observations due to the weak bonding. Nonetheless, the cross section microstructures of the oxide scales were reserved. It can be seen that these scales can be divided into two layers, i.e. a thin and dense SiO2-rich outer layer and a highly porous skeleton inner layer (the white region). During the EDS analysis of the white inner layer, Zr and O elements were detected in the grains near the outer layer while Zr and B elements were detected near the substrate. According to the XRD analysis (see Fig. 4a), this white inner layer was determined to be mainly composed of ZrO2 and ZrB2 grains. This layered microstructure is similar to that of the oxide scales formed on ZrB2-SiC composites after high temperature oxidation [28,33].  4. Discussion  4.1. Chemical reactions during the oxidation of three materials  Under the conditions used in the present work, the following oxidation reactions were expected for these three materials:  2C ðsÞ þ O2  ðgÞ ! 2CO ðgÞ  2=5ZrB2  ðcrÞ þ O2  ðgÞ ! 2=5ZrO2  ðcrÞ þ 2=5B2O3  2=3SiC ðcrÞ þ O2  ðgÞ ! 2=3SiO2  ðlÞ þ 2=3CO ðgÞ  ðR1Þ  ðR2Þ  ðR3Þ  During the oxidation of GZS at temperatures above 1600 °C, a massive amount of CO (g) was produced due to the oxidation of graphite matrix and SiC [reactions (R1) and (R3)]. Meanwhile, oxide scales formed on the surface due to the oxidation of ZrB2 and SiC in the substrate [reaction (R2) and (R3)]. These two factors could cause low oxygen partial pressures at the substrate/environment interface, which might trigger the active oxidation of SiC [reaction (R4)] in GZS. And the SiO (g) produced could escape outwards and condense back to SiO2 (l) [reaction (R5)] where the oxygen partial pressure becomes high enough.  SiC ðcrÞ þ O2  ðgÞ ! SiO ðgÞ þ CO ðgÞ  2SiO ðgÞ þ O2  ðgÞ ! 2SiO2  ðR4Þ  ðR5Þ  Besides, the sublimation of graphite and the volatilization of condensed oxides should also be considered at these high temperatures. At 1600 °C and above in air, the predominant vapour species above ZrO2 (cr), B2O3 (l) and SiO2 (l) are ZrO2 (g) [16], B2O3  Fig. 7. Surface and cross section morphologies of GZS after oxidation at (a, a0 ) 1800 and (b, b0 ) 2000 °C for 1800 s.  \\x0c', '188  Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  (g) [16] and SiO2 (g) [19], respectively. Based on the data from the online NIST-JANAF tables [34], we calculated the temperature dependence of vapour pressures of five graphite species [Cx (g), x = 1, 2, 3, 4, 5] and three oxide species [ZrO2 (g), B2O3 (g) and SiO2 (g)] according to reactions (R6-R9), as shown in Fig. 8. Assumptions of unit activity for all condensed phases and an ambient pressure of 1 atm (i.e. 1.013 \\x02 105 Pa) were used in the calculations. It can be seen that at 2000 °C, the vapour pressures of Cx (g) and ZrO2 (g) are still below \\x1810\\x008 atm and hence negligible, whereas the vapour pressures of B2O3 (g) and SiO2 (g) (PSiO2) reach 5.8 \\x02 10\\x001 and 1.6 \\x02 10\\x005 atm, respectively. The evaporation of B2O3 (l) and SiO2 (l) could cause additional weight losses and weaken the protectiveness of the oxide scale by forming bubbles and pores.  (PB2O3)  xC ðsÞ ! Cx  ðgÞ; x ¼ 1; 2; 3; 4; 5  ZrO2  ðcrÞ ! ZrO2  ðgÞ  B2O3  ðlÞ ! B2O3  ðgÞ  SiO2  ðlÞ ! SiO2  ðgÞ  ðR6Þ  ðR7Þ  ðR8Þ  ðR9Þ  4.2. Comparison of the oxidation resistance of three materials  4.2.1. Analysis of the apparent radius loss and weight loss results  For pure graphite, its oxidation proceeds with the continuous consumption of the substrate mainly by reaction (R1) at 1600 °C [35]. Because no condensed oxide phase was produced during its oxidation, both of the apparent radius loss and weight loss rates obtained truly represent its oxidation resistance. Actually, the weight loss rate of pure graphite at 1600 °C can also be derived from the corresponding apparent radius loss data. The so-derived weight loss rate, being 1.5 \\x02 10\\x004 g/cm2 s, shows close consistency with the result of the weight loss measurements (see Table 2). While for GZ and GZS, besides the oxidation of the graphite matrix, the oxidation of ZrB2 [reaction (R2)] and SiC [reaction (R3)] also take place. As a result, single-layer and adherent oxide scales form on the sample surface during oxidation at 1600 °C (see Fig. 6). In this regard, the apparent radius loss represents directly the dimensional stability of materials, which is a total radial dimension change of the sample combining the net radius loss of the substrate and the thickness of the oxide scale. However, even when taking into account of the radial increase due to the formation of the oxide scales, the composite materials still display lower net radius losses. After oxidation at 1600 °C for 1800 s, the apparent radius losses of  Fig. 8. Vapour pressure vs. temperature of five graphite species and three oxide species, calculated at ambient pressure according to reactions (R6-R9). Data are from the online NIST-JANAF tables [34].  GZ and GZS are \\x180.29 and \\x180.21 mm (Fig. 2a), while the corresponding oxide scale thicknesses are \\x18700 and \\x18225 lm (Fig. 6a0 and b0 ), respectively. It can be known that the net radius losses of GZ and GZS are \\x180.99 and \\x180.44 mm, respectively. Because of the linear characteristics of the apparent radius losses, it is reasonable that the net radius loss rates of GZ and GZS at 1600 °C are calculated to be 5.5 \\x02 10\\x004 and 2.4 \\x02 10\\x004 mm/s, respectively. It is only \\x1827% and means that the net radius loss rate of GZS \\x1844% of that of G and GZ, respectively. Therefore, the apparent radius loss results obtained by in situ dimension monitoring effectively represent the superior oxidation resistance of GZS to that of G and GZ. Similarly, the obtained weight loss rates of GZ and GZS in Table 2 also combine the net weight loss of the substrate and the weight gain of the surface oxide layer. To identify its effectiveness, we calculated theoretically the weight loss of three materials after reaction with 1 mol O2 by reactions (R1-R3), (R8) and (R9), as listed in Table 3. It can be known that, even without considering the evaporation of B2O3 (g) and SiO2 (g) [reactions (R8) and (R9)], the theoretical weight loss of GZS still reaches 75% and 91% of that of pure graphite and GZ, respectively. However, the experimental weight loss rate of GZS at 1600 °C is only \\x1824% and \\x1854% of that of G and GZ (see Table 2), respectively. It means that the lower weight loss rates of GZ and GZS have not simply resulted from the weight gain due to the production of the condensed oxides. Instead, they also truly reflect the better oxidation resistances of these two graphite-based composite materials, which are attributed to the formation of the oxide scales. Especially, compared to the improvement by single ZrB2 addition, the oxidation resistance of graphite was further enhanced greatly by the dual addition of ZrB2 and SiC.  4.2.2. Formation and effects of the porous oxide scales on two composites  Generally, the oxidation process of graphite involves the following five steps occurring in series [36]: (i) transportation of oxygen from the ambient air to the solid surface, (ii) adsorption of its molecules onto the active sites on the solid surface, (iii) reaction between the adsorbed molecules and carbon atoms at the active sites, (iv) desorption of the oxide molecules from the solid surface into the gas phase, (v) diffusion of the gaseous products through the boundary layer into the ambient air. The oxidation rate of graphite is controlled by the slowest one of these steps. Usually, steps (ii) and (iv) are considered to be extremely fast. At 800 °C and above, steps (i) and (v) are much slower than step (iii) [36,37]. In this case, the oxidation rate of graphite at 1600 °C is boundary layer diffusion-controlled. With the additions of ZrB2 and SiC, single-layer and adherent oxide scales formed on these two graphite-based composite materials during oxidation at 1600 °C (see Fig. 6). In the oxidation process, two additional steps were involved, the inward penetration of oxygen and the outward transportation of gaseous products through the oxide scales. Because massive amounts of gases, such as CO (g), B2O3 (g) and SiO (g), were produced on the substrate surface, a great number of through-thickness pores and  i.e.  Table 3 Theoretical weight losses of three materials after reaction with 1 mol O2.  Materials  After reaction with 1 mol O2  Weight loss (g)  Related reaction  G GZ  GZS  24.0  19.8(min.) 21.9(max.) 18.0(min.)  20.1  22.0(max.)  (R1) (R1), (R2) (R1), (R2), (R8) (R1), (R2), (R3) (R1), (R2), (R3), (R8) (R1), (R2), (R3), (R8), (R9)  \\x0c', 'Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  189  micro-cracks existed in the oxide scales formed on GZ and GZS. These pores and micro-cracks acted as short-path channels for the diffusion of gaseous phases. Therefore, the oxidation kinetics of GZ and GZS still followed the linear rate law. Although the oxide scales formed on these two composites were less protective, however, the surface area of the substrate exposed directly to the environment was greatly reduced due to the existence of the oxide scales. As a result, the oxidation rates of GZ and GZS became much lower than that of pure graphite. And now, the oxidation rate predominantly depends on the microstructure of the oxide scales. For the oxidation of GZ at 1600 °C, the PB2O3 on its oxide scale is as high as 9.1 \\x02 10\\x003 atm, as shown in Fig. 8. The intensive evaporation of B2O3 (l) due to its high vapour pressure resulted in a porous ZrO2 skeleton scale on the surface (see Fig. 6a and a0 ). Compared with GZ, the oxide scale formed on GZS at 1600 °C contains much amorphous SiO2 and is denser and thinner (see Fig. 6b and b0 ). This kind of 2.5 \\x02 10\\x008 of glassy SiO2 (l) has a vapour pressure atm at 1600 °C, being about five orders of magnitude lower than that of B2O3 (l) (see Fig. 8). Besides, it can seal the micro-cracks and pores in the oxide scale mainly containing ZrO2 grains [38,39]. In this way, the oxide scale formed on GZS acted more effectively as a barrier against the inward penetration of oxygen than that formed on GZ. What’s more, the adhesion of the oxide scale to the substrate becomes stronger due to the wetting of SiO2 (l). Therefore, GZS has an even better oxidation resistance than that of GZ at 1600 °C. Fig. 9 shows the relationship between the weight loss rate and the porosity of the oxide scale of three materials after oxidation at 1600 °C up to 1800 s. The oxide scale porosity of graphite was considered as 100% due to its inexistence. Good correlation can be seen to exist between these two properties.  4.3. Effect of temperature on the microstructure of the surface scale formed on GZS  Unlike the single-layer and adherent oxide scale formed at 1600 °C, the surface scales formed on GZS at 1800 and 2000 °C exhibited layered microstructures and adhered weakly to the underneath substrate (see Fig. 7). Accordingly, at these higher temperatures, the apparent radius loss curves displayed nonlinear characteristics and the weight loss rate increased. It indicates that the microstructure of the surface scale formed on GZS during oxidation is temperature-dependent and affects dominantly the oxidation rate of GZS at different temperatures. Since the oxidation rate of graphite in the boundary layer diffusion-controlled regime depends weakly on temperature [36,37], the temperature dependence of the microstructure of the surface scale results from the  change of the oxidation behaviour of ZrB2 and SiC in the graphite matrix with temperature. For that, we analyzed the thermodynamic stabilities of phases in the C-ZrB2-SiC-O2 system and then discussed the mechanism for the microstructure change of the surface scale formed on GZS with oxidation temperature.  4.3.1. Thermodynamic stability of phases in the C-ZrB2-SiC-O2 system  With the existence of ZrB2 and SiC in the graphite matrix, oxidation reactions (R1-R3) would take place during the oxidation of GZS at 1600-2000 °C in air. To understand the thermodynamic stability of phases in this C-ZrB2-SiC-O2 system, we calculated and compared the temperature dependences of the equilibrium oxygen pressure for reactions (R1-R3) [PO2(R1-R3)] based on the data from the online NIST-JANAF tables [34], as plotted in Fig. 10. In the calculations, it was also assumed that all condensed phases have unit activities and the ambient pressure is 1 atm. Since graphite is the major phase in all three materials, CO (g) becomes the predominant constituent of the gaseous species during their oxidation in air. Therefore, in the current study, it was assumed additionally that the pressure of CO (g) (PCO) is constant at a value of 1 atm. The same assumption has been made by Rezaie et al. [40] in their thermodynamic analysis of the oxidation of a 15 vol.% Ccontaining ZrB2-SiC in air at 1500 °C. It can be known from Fig. 10 that the equilibrium oxygen pressures (PO2) for these oxidation reactions increase with the increase of temperature. The value of PO2 (R2) equals that of the passive oxidation of SiC [PO2 (R3)] at \\x181456 °C, while PO2 (R1) has the same value with that of PO2 (R2) and PO2 (R3) at \\x181499 and 1512 °C, respectively. At temperatures above 1512 °C, the oxidation of C, ZrB2 and SiC exhibit the equilibrium PO2 in an increasing sequence. Therefore, different phases are stable to exist in four different PO2 regions for the C-ZrB2-SiC-O2 system above 1512 °C, as also labelled in Fig. 10. These four regions are: (1) PO2(R2) < PO2, all three oxidation reactions of (R1), (R2) and (R3) are feasible, i.e. all three components in GZS will be oxidized; (2) PO2(R3) < PO2 < PO2(R2), only reactions (R1) and (R3) are feasible, i.e. SiC and graphite matrix rather than ZrB2 in GZS will be oxidized; (3) PO2(R1) < PO2 < PO2(R3), only reaction (R1) is feasible, i.e. only graphite matrix in GZS will be oxidized while both ZrB2 and SiC stay un-oxidized; (4) PO2 < PO2(R1), none of these three oxidation reactions will take place, three components in GZS will stay un-oxidized. It should be noted that B2O3 (l) will vaporize as quickly as it is formed above 1400 °C [16,41] due to its high vapour pressures (see Fig. 8) and it is independent of PO2. Though SiO2 (l) has lower vapour pressures and is also independent of PO2, the vapour pressures of SiO (g) (PSiO) above both SiC (cr) and  i.e. all  Fig. 9. The relationship between the weight loss rate and the porosity of the oxide scale of three materials after oxidation at 1600 °C for 1800 s.  Fig. 10. Equilibrium oxygen pressure as a function of temperature for the reactions of C, ZrB2 and SiC with O2 [reactions (R1-R3)], assuming PCO = 1 atm. Data are from the online NIST-JANAF tables [34].  \\x0c', '190  Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  SiO2 (l) become considerably high at PO2 approaching PO2(R3) and rise rapidly with increasing the temperature, as shown in Fig. 11a. The PSiO at PO2(R3) and PO2(R2) (see Fig. 11b) exceeds 1 atm at temperatures above 1870 and 1988 °C, suggesting the active oxidation of SiC and the evaporation of SiO2 (l) by SiO (g) [reverse reaction of (R5)], respectively.  4.3.2. Mechanism for the microstructure change of the surface scale with temperature  It can be known from the above analysis that, at the early stage (<300 s) during the oxidation of GZS at 1600 and 1800 °C in air, the grains of C, ZrB2 and SiC on the sample surface tended to be oxidized simultaneously and formation of SiO2 (l) by the passive oxidation of SiC [reaction (R3)] was preferred due to the high PO2 in air (\\x180.2 atm). Consequently, CO (g) was continuously released and condensed oxides including ZrO2 (cr), B2O3 (l) and SiO2 (l) were also produced on the sample surface. However, the PB2O3 is already above 9.1 \\x02 10\\x003 atm at 1600 °C and reaches 9.0 \\x02 10\\x002 atm at 1800 °C (Fig. 8), and its vaporization rate increases linearly with increasing temperature [41]. Therefore, no B2O3 (l) was retained in the oxide scales of the as-oxidized GZS sample (as also in the oxide scale of the as-oxidized GZ sample). On the other hand, SiO2 (g) is the predominant vapour species above SiO2 (l) on the PO2 = \\x180.2 atm) outer surface of the oxide scale (where (see Fig. 11a). Since the PSiO2 is 2.5 \\x02 10\\x008 atm at 1600 °C and reaches only 8.6 \\x02 10\\x007 atm at 1800 °C, SiO2 (l) stayed liquid. At this early stage of the oxidation, porous oxide scales with the single-layer microstructure of ZrO2 (cr) skeleton filled with SiO2 (l) would form on GZS due to the low contents of ZrB2 and SiC additions. A boundary layer was then established with contrary potential gradients of  PSiO  PO2 = PO2  PCO and PO2 ranging from the surface of the substrate through the porous oxide scale to the ambient air. The PO2 was highest at the (\\x180.2 atm) and gradually deouter edge of the boundary layer creased across the oxide scale as the distance to the substrate decreased. When the PO2 in the oxide scale is enough low, SiO (g) becomes the predominant vapour species above all the condensed phases and the highest PSiO appears at the SiC (cr)/SiO2 (l) interface, as shown in Fig. 11a. At 1600 °C, the highest PSiO is 2.4 \\x02 10\\x002 atm [when PO2 = PO2(R3) = 1.4 \\x02 10\\x0015 atm], which is only \\x182% of its equilibrium PCO. Most of the SiO (g) escaped to the environment with CO (g) through the pores in the oxide scale. As a result, the single-layer microstructure of the porous ZrO2 (cr)/SiO2 (l) scale was sustained. At 1800 °C, however, the PSiO is already 1.6 \\x02 10\\x001 (4.2 \\x02 10\\x0013 atm), and the atm when the PO2 is down to PO2(R2) 4.2 \\x02 10\\x001 highest is as much as atm [when (R3) = 6.0 \\x02 10\\x0014 atm], which reaches \\x1840% of its equilibrium PCO. It means that a large amount of SiO (g) was released by the active oxidation of SiC (cr) [reaction (R4)] and the vaporization of SiO2 (l) [reverse reaction of (R5)]. As a result, a ZrO2 + ZrB2 transition layer were separated out from the initial ZrO2 (cr)/SiO2 (l) layer and the unaffected substrate. Because of the absence of SiO2 (l), this transition layer is more porous, which caused the exposure of more surface grains of the substrate and also weakened the adhesion of the oxide scale to the substrate. Meanwhile, a partial amount of SiO (g) condensed back to SiO2 (l) [reaction (R5)] at the outer surface of the oxide scale where the PO2 was satisfied [42,43], and the rest of SiO (g) escaped to the environment and condensed on the wall of the testing chamber. As the oxidation proceeded, more SiO2 (l) formed on the sample surface. It spread to fill the pores and form a surface glassy layer due to its lower viscosity at higher temperatures [15] and the lack of constraint from the ZrO2 (cr) skeleton [44]. In this way, however, the release path of gaseous products through the pores was also cut off and then the total gas pressure under the oxide scale gradually increased. The high gas pressure caused the bulging and cracking of the SiO2 (l)rich surface layer, which also corresponds to the fluctuations in the apparent radius loss curve shown in Fig. 2b. While at 2000 °C, the PB2O3 reaches 5.8 \\x02 10\\x001 atm, so the B2O3 (l) evaporated faster. Meanwhile, the PSiO exceeds 1 atm when the PO2 is below 2.8 \\x02 10\\x0011 atm [which is still above the 2.3 \\x02 10\\x0011 atm of PO2(R2)], and the PSiO reaches as high as 4.3 atm when PO2 decreases to 1.5 \\x02 10\\x0012 atm [PO2(R3)]. It suggests that a greatly enhanced active oxidation of SiC took place from the very beginning. As a result, a ZrO2 + ZrB2 transition layer with more ZrB2 grains were separated out, as evidenced by the XRD results that more ZrB2 was detected in the surface scale of GZS after oxidation at 2000 °C (see Fig. 4d). Besides, a much higher amount of SiO (g) was released. Therefore, it is believed that a SiO2 (l)-rich surface layer formed at a much earlier time during oxidation (before 60 s). Since SiO2 (l) has an even lower viscosity at 2000 °C [15], the high pressure gaseous products under the oxide scale caused the bulging of the SiO2 (l)-rich surface layer and broke through it more easily, as evidenced by the surface morphology in Fig. 7b that more bulges and pores existed in the surface layer at 2000 °C. This property helped in releasing the high gas pressure and prevented the catastrophic disruption of the oxide scale. However, it is conceivable that the weak adhesion of the surface scale to the substrate and the accumulation of gaseous products beneath the surface scale will eventually result in the detachment of the surface scale.  Fig. 11. (a) Volatility diagram for SiC at 1600, 1800 and 2000 °C, and (b) vapour pressure of SiO (g) vs. temperature at PO2(R1, R2, R3), assuming PCO = 1 atm. Data are from the online NIST-JANAF tables [34].  Using an induction-heating based ultra-high temperature oxidation testing apparatus, the effects of ZrB2 and ZrB2 + SiC additions on the isothermal oxidation behaviour of graphite at  5. Conclusions  \\x0c', 'Z.-H. Gao et al. / Corrosion Science 76 (2013) 182-191  191  1600-2000 °C in air were investigated. The oxidation kinetics were characterized by the in situ dimensional monitoring as well as the ex situ weight loss measurement. The major conclusions are presented below.  1. The oxidation of pure graphite, C-ZrB2 and C-ZrB2-SiC at 1600 °C in air all follow the linear rate law. However, C-ZrB2- SiC exhibits the lowest apparent radius loss rate of 1.0 \\x02 10\\x004 mm/s, which is an order of magnitude lower than that of pure graphite and 63% of that of C-ZrB2. 2. Compared with the single addition of ZrB2, the dual addition of ZrB2 + SiC improves the oxidation resistance of graphite at 1600 °C more effectively, because the oxide scale formed during oxidation is denser and has a stronger adhesion to the substrate due to the existence of glassy SiO2. 3. The oxidation rate of C-ZrB2-SiC at 1600-2000 °C is dominated by the microstructure of the surface scale formed on its surface. The greatly enhanced active oxidation of SiC contributes to the formation of the layered microstructure of the surface scales and the increased oxidation rates at 1800 and 2000 °C. 4. The thermodynamic analysis indicates that the addition of SiC may no longer be effective on improving the oxidation resistance of graphite above \\x181870 °C in that the pressure of SiO (g) produced by the active oxidation of SiC in the graphite matrix will exceed 1 atm and eventually cause the detachment of the oxide scale.  Acknowledgements  The authors would like to thank Prof. H. -M. 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Hu, X. Zhang, Y. Liu, W. Han, Effects of oxygen partial pressure and atomic oxygen on the microstructure of oxide scale of ZrB2-SiC composites at 1500 °C, Corros. Sci. 73 (2013) 44-53. J. Li, T.J. Lenosky, C.J. Först, S. Yip, Thermochemical and mechanical stabilities of the oxide scale of ZrB2 + SiC and oxygen transport mechanisms, J. Am. Ceram. Soc. 91 (2008) 1475-1480.  carbon combustion in  [30]  [44]  \\x0c']"
},{
  "_id": 67,
  "PDF": "Embedded ZrC-SiC nanocomposites from hydrothermal precursor with temperature-dependent oxidation resistance and high sinterability.pdf",
  "Text": "['Journal of Alloys and Compounds 791 (2019) 316e327  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : h t t p : / / w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  Embedded ZrC-SiC nanocomposites from hydrothermal precursor with temperature-dependent oxidation resistance and high sinterability  Wentao Xu a, Youfu Zhou a, *, He Lin a, Shuai Lu a, b, Zhiguang Wang c, Kun Wang d  Junrong Ling a, b, Rui Wang a, b,  a Key Laboratory of Optoelectronic Materials Chemistry and Physics, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences,  Fuzhou 350002, China b University of the Chinese Academy of Sciences, Beijing 100049, China c Institute of Modern Physics, Chinese Academy of Science, Lanzhou 73000, China d School of Materials Science and Energy Engineering, Foshan University, Foshan 528000, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 3 December 2018 Accepted 21 March 2019 Available online 23 March 2019  Keywords:  ZrC-SiC nanocomposite Hydrothermal precursor Oxidation resistance Embedded morphology Spark plasma sintering  ZrC-SiC nanocomposites were synthesized via a facile, green hydrothermal precursor conversion method. The zirconium, silicon and decomposed glucose can be successfully incorporated into a homogeneous framework. In subsequent pyrolysis, the precursor transforms to a core-matrix structure of the nanoZrO2 and amorphous carbon/silica mixture, resulting in a short diffusion path and limited grain growth. Carbide first appears at a low temperature of 1200 \\x0e C, and the conversion is completed at 1500 \\x0e C with low oxygen content. The composite consists of grains about 100 nm, exhibiting specific embedded morphology, and has different oxidation resistance in three temperature zones based on the component. The origin mechanism and properties have been elucidated and analyzed. The present work demonstrates the effectiveness of hydrothermal chemistry for the synthesis of carbide composites and their promising application in high temperature protection. Such nanocomposites with controllable morphology and high sinterability beneficial for subsequent densification are also verified. The additivefree ceramic has been obtained with almost fully density (relative density >99%) at a low temperature of 1700 \\x0e C by spark plasma sintering (SPS). The sintered specimen possesses fine microstructure with hierarchical grain size distribution (about 1 mm and 100 nm, respectively) and good mechanical properties (fracture toughness of 4.3 ± 0.4 MPa m1/2 and Vickers hardness of 22.8 ± 0.7 GPa). © 2019 Published by Elsevier B.V.  1.  Instruction  Zirconium carbide (ZrC) as one of high-potential ultra-high temperature ceramics (UHTCs) has the inherent high melting point, high hardness, good wear and corrosion resistance. It has a promising application in the aerospace field, such as rocket engine and reentry heat shields of hypersonic vehicles [1e3]. In addition, due to the neutron transparency and weak damage sensitivity under irradiation, ZrC is also considered as a hopeful candidate for fission barrier and inert matrix fuel in nuclear industry [4,5]. However, the poor oxidation resistance under extreme conditions limits its  * Corresponding author. E-mail address: yfzhou@fjirsm.ac.cn (Y. Zhou).  https://doi.org/10.1016/j.jallcom.2019.03.299 0925-8388/© 2019 Published by Elsevier B.V.  widespread use. In order to overcome this obstacle, silicon carbide (SiC) is introduced, which can form a dense SiO2 protective layer at an elevated temperature [6,7]. Based on the second phase particle toughening, ZrC-SiC composites also have enhanced mechanical strength, especially for nanograin materials. So it has been expected that such carbide composite presents better comprehensive properties, compared with single-phase ceramics [8e10]. The ZrC-SiC composites can be simply prepared by solid blend methods, including reduction of oxides, or ball-milled multicomponent carbides [11e13]. However, these rough routes are energy and time consuming, and the final product is agglomerated and contaminated because the ingredients mix in a coarse scale and the abrasion of the grinding media increases the impurities. Solution-based alternatives are proposed, such as polymer-derived precursors and sol-gel conversion methods, which have the  \\x0c', 'W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  317  advantage of homogeneous distribution at the molecular level and conversion achieved at low temperature. For the polymer-derived precursor method, soluble zirconium(zirconium alkoxides/polyzirconoxane (PZO), [(C4H8O)Zr(acac)2]n (PZC)) and silicon(polycarbosilane (PCS), polymethylsilylacetylene (PMSA)) precursors were blended and cured to form preceramic polymers [14e16]. For the sol-gel method, Zr-, Si-sol was prepared and mixed from related salts or alkoxides, and then cross-linking, aging and drying was carried out to form xerogel [17e19]. All the precursors and xerogel were subjected to pyrolysis for ZrC-SiC composites under 1400e1700 \\x0eC. Nevertheless, shortcomings of existing solutions are still not to be ignored, for example, the loss of SiC component is a common problem due to the evaporation of gaseous intermediate. Beyond that, the critical synthesis conditions, complex operations, toxicity of raw materials, high cost and technical safety restrict the wide application of above methods. Therefore, it is still a challenge to produce high-quality ZrC-SiC nanocomposites through an industrially scalable process. In recent decades, the hydrothermal strategy that triggers chemical reaction and change in solubility in sealed heating liquid above ambient temperature and pressure is considered to be a convenient and green way to prepare nano/ micro materials [20,21]. However, such methods for synthesizing non-oxides are rare compared with oxides [22]. In our previous works, we have successfully synthesized AlN and ZrC nanopowders from hydrothermally derived precursors without the need of air or moisture sensitive precautions [23,24]. It is worth mentioning that, due to strong covalent bonds and low diffusion coefficients, dense ZiC-SiC body is hard to obtain, which limits its engineering applications [25,26]. Most of the previous works focused on reactive hot pressing (R-HP) or spark plasma sintering (R-SPS) process, and it is hoped that the reaction driving force can lower the sintering temperature (the reaction system includes: ZiSi2-C, ZrSiO4-C, ZrCSi-graphite, ZrC-Si, ZrC-polycarbosilane and SiC-ZrH2) [9,25,27e31]. However, the study on microstructure control and sinterability improving by regulating the particle size and morphology of starting nanocomposites, is still insufficient [13,32]. Present work attempts to establish a low-cost and convenient method for the preparation of high-quality ZrC-SiC composite. It also provides an opportunity to investigate relationship between the microstructure of ceramics and initial nanocomposite. The glucose (Glu) with high carbon conversion yield and hydroxyl groups is introduced as carbon source and polydentate ligand, which can chelate zirconium and silicon at the atomic/molecular level by hydrothermal treatment. The Si/Zr-O-C bonds in the homogeneous precursor can reduce the mass transfer distance in the carbothermal process, which is beneficial for the reaction thermodynamics and kinetics. This method is versatile and can be extended to other carbide composites. The characters, morphology evolution and oxidation resistance of the ZrC-SiC nanocomposites  have been studied in detail. In addition, the solidification via SPS, which can induces particle activation and local melting in a short time suppressing grain growth, has been carried out. The densfication of the sintered specimen and the effect of composite on microstructure and mechanical properties were discussed.  2.  Experimental section  2.1. Materials  Zirconium oxychloride octahydrate (ZrOCl2$8H2O, ZOC, 99%) and glucose (Glu, 99%) were purchased from Aladdin Chemical Co., China, used as received. Tetraethoxysilane (TEOS, 98%) and ethanol (99.7% grade) was purchased from Adamas Chemical Co., China, used as received. Deionized water (DI water) was prepared inhouse by using a Thermo Scientific Barnstead Easypure II system.  2.2.  Preparation of precursors and ZrC-SiC composites  In a typical synthesis, ZOC and TEOS (a total of 0.06 mol) were dissolved in 50 mL of DI water, and then an appropriate amount of Glu was added to adjust Glu/(Zr þ Si) molar ratio with continue stirring for 5 h at room temperature. The Zr/Si molar ratio in the final composite was controlled at 2, 1 and 0.5, respectively. The clear solution was sealed in a 100 mL Teflon-lined stainless steel autoclave at 180 \\x0eC under autogenous pressure for 20 h. After cooling to room temperature at a rate of 5 \\x0e C/min, the resulting brown paste was filtered, washed with DI water and ethanol to the pH of 7 and vacuum dried at 80 \\x0e C. These precursors with different Zr/Si ratios are denoted as P-ZS21, P-ZS11 and P-ZS12, respectively. All samples were heated in a tube furnace with a heating rate of 5 \\x0eC/min under a flowing argon atmosphere from 900 \\x0e C to 1600 \\x0eC for 1 h (denoted as ZS21, ZS11 and ZS12). The pyrolysis product can be easily crushed into powder in a mortar. The illustration of the synthesis procedure is shown in Fig. 1. As a comparison, ZrC-SiC composite was also prepared from the precursor by the mixing/ drying protocol instead of the hydrothermal reaction (denoted as ZS11mix) as follows: the appropriate amount of Glu was dissolved in the ZOC/TEOS solution, stirred for 5 h, and the solution was removed by rotary evaporation, obtaining brown powder. The dried precursor is then converted to carbide by similar carbothermal reduction.  2.3. Densification of ZrC-SiC composites  Such synthesized ZrC-SiC powder without further treatment and free of sintering additive was poured into a graphite die (inner diameter of 10 mm), and sintered using SPS equipment (SPS-211Lx, Fujidempa, Japan) in vacuum. The temperature, measured by a vertically located optical pyrometer through a hole in the upper  Fig. 1.  Illustration of the synthesis procedure of ZrC-SiC composites.  \\x0c', \"318  W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  graphite punch, was increased to 1600 \\x0e C at 100 \\x0eC/min up, and then at 50 \\x0e C/min up to 1700 \\x0eC for 10 min henceforth. After that the power was switched off for fast cooling. A uniaxial pressure of 45 MPa was loaded and kept constant until the end of the dwell period.  2.4.  Characterization  Fourier transformation infrared spectra (FTIR, Spectrum One, PerkinElmer, USA) of starting reagents and precursors were recorded within the range of 4000\\x00400 cm \\x001. The phase compositions were examined by X-ray diffraction (XRD, Ultima-IV, Rigaku, Japan) using Cu Ka radiation. Then the lattice parameters of ZrC phase were determined by indexing and least-squares refinement with the MDI Jade 5 software [33]. The morphology and microstructure of the products were characterized by field emission scanning electron microscope (FESEM, SU-8010, Hitachi, Japan) equipped with an energy dispersive spectroscopy (EDS) apparatus, and transmission electron microscopy (TEM, Tecnai F20, FEI, USA).  Fig. 2.  FTIR spectra of starting reagents and precursor.  Thermogravimetric analysis (TGA, STA 449F3, Netzsch, Germany) for precursors and composites was performed in argon or air atmosphere from room temperature to 1500 \\x0e C respectively, at a heating rate of 10 \\x0e C/min. Oxygen content was detected by an oxygen analyzer (ONH836, LECO, USA). The SPS sample was ground and polished with a diamond paste to 0.5 mm and chemically etched using a mixed acid (HF: HNO3: H2O, volume ratio ¼ 1:1:2) for microstructure investigation. The bulk density was measured by the Archimedes method. Vickers hardness (HV) was measured under loads of 9.8 N for 15s on a polished surface. Fracture toughness (KIC) was evaluated by the direct crack measurement (DCM) method, using the equation of Niihara [34], suitable for the ratio of cracks length to diagonal length larger than 2.5. In this equation, Young's modulus of composite was calculated by a mixture rule [35]. More than 15 measurements were performed to get the average value of hardness and toughness.  3. Results and discussion  3.1.  Characterization of  the hydrothermal-derived precursors  In the internal environment under hydrothermal condition, the reactants (ZOC, TEOS and Glu) undergo multiple reactions to form polymeric precursor. FTIR analysis is used to verify the conversion. Fig. 2 shows the profile of the precursor (P-ZS11), where several peaks change significantly compared to raw materials. The main \\x001) band attributed to the stretching vibration of eOH (Glu 3410 cm \\x001, was partially shifted to a lower wave number of 3201 cm indicating that the eOH group was involved in the bonding [36]. It is worth mentioning that the peak from the vibration of eCHO (Glu \\x001) disappears, with the appearance of a carboxyl (eCOO) 2940 cm \\x001 and 1396 cm \\x001 respectively. and Zr-O-C absorption at 1703 cm This can be attributed to the oxidation and partial carbonization of Glu, forming humin in subcritical water [37]. The absence of char\\x001) and simultaacteristic peaks of Si-O-C (1103, 965 and 793 cm \\x001), \\x001) neous build of Si-O-Zr (1047 cm Zr-O (656 and 459 cm indicate the hydrolysis and subsequent condensation for both TEOS and ZOC [38,39]. Thus, zirconium, silicon and decomposed Glu have been successfully incorporated into the framework of complex. With this in mind, the possible reactions during precursor formation are the hydrolysis and partial condensation of ZOC and TEOS, accompanied by binding with decomposed Glu (Humin) (Fig. 3).  Fig. 3. The possible reactions during hydrothermal preparation.  \\x0c\", 'W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  319  3.2.  Preparation of  the ZrC-SiC composites  Preparation of ZrC or SiC from an oxide-carbon precursor generally represented by the following equations:  is  ZrO2 (s) þ 3C (s) ¼ ZrC (s) þ 2CO (g)  SiO2 (s) þ 3C (s) ¼ SiC (s) þ 2CO (g)  (1)  (2)  Eq. (1) or (2) is a simplified carbothermal reduction that actually consists of a series of solid-solid, solid-gas, and gas-gas reactions [11,40e42]. It is necessary to determine and optimize the C/(Zr þ Si) carbon. The precursors with different Glu/(Zr þ Si) molar ratio (0.7, molar ratio to produce a stoichiometric ZrC-SiC without residual 0.8 and 0.9) were prepared and then pyrolyzed at 900 \\x0eC under Ar flow for 1 h (denoted as ZS11-07, ZS11-08 and ZS11-09). The decomposed precursors containing ZrO2, SiO2 and carbon were  Fig. 4. TG-DTG-DTA curves of precursor (P-ZS11) in Ar atmosphere.  then heated and quantified by the mass change in air. Fig. S1 shows the typical thermogravimetric (TG) curves. The precursors have carbon contents of 24.07, 28.15, and 30.8 wt%, which correspond to C/(Zr þ Si) ratios of 2.42, 2.99, and 3.23, respectively. According to above equations, the composition of the final precursor was adjusted to about 3:1 M ratio of C/(Zr þ Si) for the subsequent process. Fig. 4 shows the TG-DTG-DTA curve of precursor (P-ZS11) as representative. Due to the absorption of moisture, the TG curve shows a rapid but slight weight loss below 120 \\x0eC. The significant loss between 120 and 500 \\x0eC can be weight attributed to the decomposition and carbonization of Glu-derived humin, resulting in an endothermic peak at 387 \\x0eC. In the stage of 500e1210 \\x0eC, the weight reduction (8.2%) is probably due to further carbonization and zirconia crystallization. Above 1210 \\x0e C, the mass loss accelerates again until 1450 \\x0eC, where the DTG curve shows a maximum peak around 1400\\x0eC corresponding to the carbothermal reduction. The pyrolyzed yield was about 40.2 wt%. The phase composition of the pyrolysis products at different temperatures was characterized by XRD analysis. As shown in Fig. 5, monoclinic (m-ZrO2) and tetragonal zirconia (t-ZrO2) exist simultaneously from 1200 to 1400 \\x0e C, because the crystallite sizes of two phases are similar and closed to the critical value [43]. A small amount of SiC (JCPDS card 73-1665, cubic) appeared at 1200 \\x0eC, while the silicon and carbon component remains amorphous. With the increase of temperature to 1400 \\x0eC, the peaks of mZrO2 and t-ZrO2 gradually decrease. SiC can be formed at an earlier (1400 \\x0e C). stage, compared with ZrC (JCPDS card 35-0784, cubic) This means that the carbothermal reduction of the silicon component occurs more easily under the same condition. Above 1500 \\x0e C, the products are carbides without any oxide impurities. This is a low synthetic temperature compared with the traditional solid state synthesis [11,40], which confirms the high reactivity of such hydrothermal precursor. For samples obtained from 1400 to 1600 \\x0eC, using Si as external standard, ZrC-related peaks were observed to shift slightly to low  Fig. 5. XRD patterns of the precursor (ZS11) treated at different temperatures, and zoom of the framed area showing the slight shift in 2q of (2 2 0) peak.  \\x0c', '320  W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  Table 1 The main characteristics of ZrC-SiC composites at different temperatures.  sample  phase  m-ZrO2 content (wt% of ZrO2)  SiC content (wt% of carbide)  ZrC lattice parameter (Å)  ZrC stoichiometry  Theoretical value  Experimental value  Oxygen content (wt%)  ZS111200 ZS111300 ZS111400 ZS111500 ZS111600 ZS211600 ZS121600  t-ZrO2, m-ZrO2, SiC t-ZrO2, m-ZrO2, SiC t-ZrO2, m-ZrO2, ZrC, SiC ZrC, SiC ZrC, SiC ZrC, SiC ZrC, SiC  32.1 54.4 35.0 / / / /  / / 28.0 28.0 28.0 16.2 43.7  / / 36.8 28.5 27.9 16.0 44.1  / / 4.691 4.694 4.697 4.695 4.695  / /  ZrC0.89O0.09 44 45 ZrC0.94O0.05 ZrC0.96O0.02 45 ZrC0.94O0.05 ZrC0.94O0.05  / / / 1.85 0.86 1.35 0.83  angles (Fig. 5). This phenomenon implies the presence of carbon oxides (ZrCxOy), in which the carbon atoms (atomic radius 0.77 Å) are partly replaced by oxygen atoms (atomic radius 0.64 Å). The composition can be estimated using lattice parameters inversely proportional to the oxygen content, based on the refinement of XRD data and the evolution of the C/O ratios detected in different oxycarbides previously [44,45]. As shown in Table 1, the phase obtained at 1600 \\x0e C with the lattice parameter of 4.697 Å, indicates a near complete transformation with little oxygen remaining. The related lattice parameter reduces as temperature decrease. However the value of SiC varies very little with temperature (from 4.360 to 4.358 Å), indicating that SiC has been completely converted at low temperature. Fig. 6 shows the XRD patterns of samples with different ZrC/SiC ratios (ZS21, ZS11 and ZS12) pyrolyzed at 1600 \\x0eC. The peak of ZrC dominates, considering the weak X-ray diffraction of SiC. It is known that silicon monoxide (SiO) gas escapes during the reaction, so silicon loss was very common in the previous reports [14e17]. Current work is also instructive in ensuring preset ratio, achieving near-theoretical ZrC-SiC composition (calculated by the RIR method and listed in Table 1) with very small Si content loss. This property and the above-mentioned high reactivity may be related to the specific microstructure of the hydrothermal precursor. As a comparison, the sample prepared from the nonhydrothermally treated precursor (ZS11mix) had much lower SiC content (18.2 wt%), higher oxygen content (3.15 wt%) and phase  Fig. 6. XRD patterns of ZS21, ZS11 and ZS12 heated at 1600 \\x0e C.  separation appearance (Fig. S2). This can be attributed to heterogeneously physical mixing and long diffusion path. The FTIR spectra of the precursor (ZS11) calcined at 1200e1700 \\x0e C were shown in Fig. S3. The peaks for Zr/Si-O bonds \\x001) observed from 1200 to 1400 \\x0e C has van(521, 739, and 813 cm ished with temperature increasing, while the peaks assigned to Zr\\x001) and Si-C (1388 cm \\x001) appear and strengthen from C (842 cm 1500 to 1700 \\x0e C [7,18]. It can be regarded as the transformation of oxide to carbides, which is consistent with the XRD results.  3.3. Microstructure and mechanism analyses  Fig. 7 shows the SEM micrographs of ZS11 obtained from 1400 to 1600 \\x0e C. In the secondary electrons (SE) mode (left side: Fig. 7a, c and e), the grains present a homogeneous near-spherical morphology at low temperature, while a few floes are considered as free carbon. As the temperature rises, the residual oxides transform into carbides gradually and the surrounding carbon is continuously consumed. The particles exhibit a polyhedral morphology with the sizes increasing from tens to several hundred nanometers. Based on the images from back-scattered electrons (BSE) (right side: Fig. 7b, d and f), the component distribution can be clearly identified according to the atomic number (Z). Bright particles with high Z-value and dark ones with low Z-value correspond to the Zr-substance and the Si-substance, respectively. The two components intermix homogeneously, exhibiting special embedded morphology, which is rarely reported [46]. A typical pattern can be contrasted by the particles circled in Fig. 7e and f, where the seemingly monolithic particles consist of two components (ZrC and SiC) in the mosaic model. When the SiC content changes (ZS21, ZS12), there is no obvious difference in morphology and particle size except the degree of aggregation, as shown in Fig. S4. The direct explanation is that the surrounding homogeneous matrix is an effective “shell” that will control mass transfer and prevent the mutual contact. The grain growth rate does not change significantly, while the long-lived SiC nanograins are more likely to aggregate. In addition, the EDS mapping of ZS11 is carried out to confirm the uniform element distribution of Zr, Si, and C (Fig. S5). TEM analysis further indicates the crystallinity and embedded distribution of ZrC-SiC composites. As shown in Fig. 8a, the particles in the filmy amorphous matrix can be observed. The grain size is in the range of 30e100 nm, which is in good agreement with the calculated results. The selected area electron diffraction (SAED) pattern of the entire region shows the continuous diffraction rings corresponding to the (111), (200), (220) lattice planes of ZrC and (111) of SiC, respectively (the orange and blue circles in Fig. 8b). The high-resolution transmission electron microscopy (HRTEM) image  \\x0c', 'W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  321  Fig. 7. SEM micrographs of ZS11 heated at various temperatures in SE (left) and BSE (right) observation modes: (a) and (b) for 1400 \\x0e C, (c) and (d) for 1500 \\x0e C, (e) and (f) for 1600 \\x0e C.  of the selected region (Fig. 8c)), confirms the embedded morphology, where the lattice fringes in an intact grain are respectively consistent with the lattice spacing of ZrC (111) (0.27 nm) and that of SiC (200) (0.21 nm) (blue rectangular area in Fig. 8a). A single ZrC crystallite (an orange rectangular area in Fig. 8a) was also observed (Fig. 8d). The associated fast Fourier transform (FFT) is shown as inset. All the crystallites are covered by continuous amorphous edge, with a thickness of 2e3 nm, which is consistent with the results in literature [23,47]. Such region can be considered as the mass transfer area. In order to explore the reaction in detail, the precursor (ZS11) was treated in Ar at 900 \\x0eC. For this early phase, loosely agglomerated sponge-like network (consisting of Zr, Si, O and C) can be observed with the XRD peaks of t-ZrO2 (Fig. 9a). HRTEM image depicts that the nanoparticles embedded in the amorphous matrix was less than 10 nm (Fig. 9b). The SAED pattern also indexes the characteristic (101), (112) and (211) lattice spaces of t-ZrO2. During the heating, crystalline zirconia forms first, while the carbon and silicon components are still amorphous. Such matrix can not only reduce agglomeration of zirconia, limit grain growth, but also increase the contact area between ZrO2 and carbon. Naturally, the carbon-silicon matrix with uniform element distribution and ultrashort diffusion path exhibits high reactivity. So the carbothermal  reduction can occur at a low temperature (about 1200 \\x0e C). Considering the formation of carbides at different temperatures, we can deduce the different reaction mechanisms for SiC and ZrC, respectively, which can be utilized to elucidate the origin of the embedded morphology. The overall carbothermal reduction of silica proceeds via formation of SiO(g), CO(g) and CO2(g) as intermediates (Eqs. (3)e(5)) [41,42,48].  SiO2 (s) þ C (s) ¼ SiO (g) þ CO (g)  SiO (g) þ 2C (s) ¼ SiC(s) þ CO (g)  SiO (g) þ 3CO (g) ¼ SiC (s) þ 2CO2 (g)  (3)  (4)  (5)  In such system, Eq. (3) as the initial step occurs at the low temperature, requiring a very low partial pressure of CO. This facilitates the subsequent Eq. (4) but vice versa for Eq. (5). Therefore, Eqs. (3) and (4) should play a dominant role at the beginning of carbothermal reduction. Since the reactants of Eq. (3) are all solid phases, the mixing uniformity and size of the SiO2 and C components in the precursor are critical. As the reaction proceeds, the partial pressure of CO increases and Eq. (5) cannot be ignored. The gaseous species (SiO(g), CO(g)) adhere to the carbon-rich matrix,  \\x0c', '322  W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  Fig. 8. TEM images of the ZS11: (a) overview, (b) selected area electron diffraction (SAED) pattern of the particles, (c) and (d) enlargement of the selected region and its associated fast Fourier transform (FFT) pattern.  (a) TEM image of precursor (ZS11) treated at 900 \\x0e C under low magnification (the insets show the XRD pattern and EDS elemental analysis), (b) HRTEM image of embed ZrO2 Fig. 9. crystallites (the inset shows SAED pattern).  giving the SiC nuclei. The succeeding growth is regulated by the external amorphous layer (see in Fig. 8c). According to the thermodynamic analysis (Fig. S6), Eq. (3) occurs above 1789 K at a low CO partial pressure (10 kPa), and the reaction temperature increases to 1994K at a high CO partial pressure (100 kPa). Eq. (4) is less affected by temperature and gas  partial pressure, which can occur in the whole temperature range free energy (about \\x0077 kJ mol \\x001). Eq. with a stable Gibbs (5) is prone to occur under lower temperature (1221 K) with high CO partial pressure (150 kPa). It is worth noting that the calculation is based on conventional synthesis. Considering the as-prepared precursor has high reactivity, the boundary of the reaction  \\x0c', 'W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  323  particle with respective epitaxial carbon tends to agglomerate for more carbon contact spontaneously. When the nearby carbon is depleted, the size of agglomerate no longer increases and carbon monoxide promotes the further reduction. The oxycarbide maturates gradually, involving absorption of CO at the external layer, inwards diffusion of carbon and outwards diffusion of oxygen. The disordered surface region has also emerged to regulate the mass transfer between the gas medium and the oxycarbide. During the formation of the Zr-containing agglomerates, the existing SiC particles will be wrapped in. And then the grains grow and intercalate along each other, forming the embedded morphology. The schematic representation can be seen in Fig. 11. It is worth mentioning that the gas phase generated during the progress, such as SiO, CO, etc., can be captured by the amorphous spongy matrix and does not escape easily. So the completion of the reaction without the component deviation can be guaranteed.  Fig. 10. The grain sizes of ZS11 composite obtained by H-W method analysis. Standard deviations are shown.  conditions can be further extended reasonably. On the other hand, the solid state reaction through the aggregation mode rather than simply centripetal diffusion of carbon into the oxide has been proposed for the formation of ZrC, according to following half-reactions [40,47]:  ZrO2 (s) þ C (s) / ZrOxCy (s) þ CO (g)  ZrOxCy (s) þ C (s) / ZrC (s) þ CO (g)  ZrOxCy (s) þ CO (g) / ZrC (s) þ CO2 (g)  (6)  (7)  (8)  As shown in Fig. 10, it can be seen that oxide and carbide sizes calculated by the Halder-Wagner (H-W) method is not identical [49]. When the precursor is heated, the ZrO2 grains gradually grow to 30e40 nm overlapping with 120 nm of ZrOxCy, and the oxycarbide size has almost no change once formed. A reasonable explanation is that more than two neighboring zirconia particles contract, sinter and then form carbides. The driving force may come from the fact that the oxides during carbothermal reduction need to consume the surrounding carbon continuously. So each oxide  3.4. Oxidation behavior of ZrC-SiC compositions  In order to study the oxidation resistance, ZrC-SiC composites were evaluated by the value of oxidation advancement (x), which was defined as the ratio of experimental weight gain to theoretical weight gain. The “theoretical weight gain” is inferred from the complete oxidation of carbide with selected composition. Each sample was tested by TG analyses under airflow from 30 to 1500 \\x0eC (Fig. 12). In the first section, the oxidation initiates at about 310 \\x0e C, till 630 \\x0eC, with a maximum peak at about and then levels off 530 \\x0eC (Fig. 12a). The exothermic peaks and weight gain/loss peaks are observed for the DTA and DTG curves, respectively (Fig. 12b and c). In this range, the mass gain can be attributed to the dominated oxidation of ZrC and few of SiC with formation of ZrO2, SiO2 and carbon. Then the gas evolution by oxidation of free carbon results in the following mass loss (Eq. (9), (10)). As the temperature rises, the x value remained stable until about 1200 \\x0e C, where the oxidation of SiC accelerates (Eq. (11)). The DTA and DTG curves are also in good agreement with that, showing the corresponding weight gain and exothermic peaks at about 1400 \\x0eC.  ZrC (s) þ O2 (g) / ZrO2 (s) þ C (s)  C (s) þ O2 (g) / CO (g)/CO2 (g)  SiC (s) þ O2 (g) / SiO2 (s) þ CO (g)/CO2 (g)  (9)  (10)  (11)  Fig. 11. Representation of the formation mechanism and composition evolution of ZrC-SiC composite from 900 \\x0e C to 1600 \\x0e C.  \\x0c', '324  W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  (<506 \\x0e C, denoted as stage I), sample ZS11 with proper ZrC content, showing the best antioxidation (the lowest x value), because ZrC is oxidized to form zirconia, which has some low temperature protective effect on the remaining carbides. In the medium temperature range (506e1365 \\x0eC, denoted as stage II), SiC is oxidized to form a dense silica protective layer on the surface, which plays a role in antioxidation. According to the data at 1200 \\x0e C as major representatives (green spots in Fig. 12a), the oxidation advancement for ZS21, ZS11 and ZS12 is only 52.4%, 51.3% and 36.4%, respectively. So the oxidation resistance significantly improves with the increased SiC content. In the high temperature region (>1365 \\x0e C, >1411 \\x0e C, especially denoted as stage III), SiC is completely oxidized and then melts or volatilizes without protection. In contrast, zirconia component is still stable, which can serve as an oxygen diffusion barrier providing some protection. So samples with high ZrC content are favorable for better oxidation resistance. Based on above results, we can select the composite with appropriate components according to different working temperature requirements, to achieve the best antioxidant effect in practical applications. Such temperature-dependent oxidation resistance may also be related to the special morphology and some synergistic effects of the composites. This hypothesis is still under investigation.  3.5.  Sintering of as-synthesized ZrC-SiC nanocomposite  It is well known that the size and uniformity of starting particles is related to the densification kinetics, microstructure and properties of ceramics [50,51]. In order to inspect the sintering activity, SPS process for ZS12 composite at a relative low temperature of 1700 \\x0e C was carried out as a representative. As shown in Fig. 13, the 1086 \\x0e C, thermal dilatation was distinct up to and then the shrinkage of the specimen initiated at 1189 \\x0eC and the densification rate increased sharply with a maximum value at 1383 \\x0eC. Finally, the shrinkage of the compact piece is almost completed at 1700 \\x0eC with the beginning of isothermal stage. According to the XRD pattern and the element mapping of the selected zone (Figs. S7 and S8), the phase composition of sintered Z12 ceramic is in good agreement with that of ZS12 powder without deviation. The BSE images of fracture surface for the sintered specimen and its EDS analysis is shown in Fig. 14a. The composite possesses a fine  Fig. 12. (a) Oxidation advancement (x), (b) DTG and (c) DTA curves of as-synthesized ZS21, ZS11, ZS12.  It is interesting to find the oxidation resistance of composite does not vary linearly with the SiC content [15]. As shown in Fig. 12a, the x values change in different temperature ranges, indicating the temperature-dependent property. At low temperature  Fig. 13. The displacement 1700 \\x0e C.  curve of ZS12 composite during the SPS experiment at  \\x0c', 'W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  325  appear within the micron domains as inclusions (marked by orange and blue arrows, respectively), forming the intragranular structure. Fig. 14b shows the micrographs of polished surface, where two phases locate along and within each other in line with the fracture surface. The average size of the micron grain is 1.08 and 1.06 mm for ZrC and SiC, respectively. While the value for the tiny spherical inclusion is 90 and 105 nm, respectively. The intergranular particles grew about 10 times larger during sintering, while the intragranular ones hardly grew compared with the starting powders. Obviously, these nanoparticles are wrapped by the other phase, preventing the further growth and aggregation. So the embedded morphology of the original nanocomposite can be inherited as expected. Compared with pure ZrC [24], the extraordinary grain growth has been inhibited due to the pinning effect [52]. By prolonging etching time to dissolve ZrC, the distinct microstructure can be exposed, where the hierarchical size distribution is further confirmed (Fig. S9). The sample exhibits multiple fracture modes including inter-, intra-granular fracture and grain pull out (the pits and outcrops) (Fig. 14a). Meanwhile the folds related to plastic grooving, indicating the some plasticity in this composite. So the improved indentation fracture toughness (KIC) of 4.3 ± 0.4 MPa m1/2 has been achieved. The specimen is fully dense (4.55 ± 0.06 g/cm3, >99%), with Vickers microhardness (HV) of 22.8 ± 0.7 GPa. Such values are even better than some products prepared by SPS under higher temperature and pressure conditions (shown in Table 2), confirming the excellent sinterability of as-synthesized nanocomposite.  4.  Conclusion  In present work, a simple one-pot hydrothermal method for synthesizing ZrC-SiC nanocomposites is proposed. As a general approach, it can be extended to the other carbide composites conveniently. The obtained hydrothermal precursor has a corematrix structure that combines the advantages of short diffusion distances and homogeneousness. It can be converted to nanosized products at low temperature and avoid compositional deviation SiC and ZrC begin to form at 1200 \\x0e C and 1400 \\x0e C, effectively. respectively, and the reaction is completed at 1500 \\x0eC. The grains exhibit an embedded morphology with the size about 100 nm, which can be explained by multiple reaction mechanisms. The understanding of the reaction process is a key element of future research. Studies on the oxidation behavior show that the composite has temperature-dependent oxidation resistance. Zirconia and silica play a major protective role in three temperature ranges (<506 \\x0e C, 506e1365 \\x0e C and >1365 \\x0e C), respectively, which can be used to guide the selection of composite under different environments. Such nanocomposite has high sinterability. The fully dense ceramic (relative density >99%) could be achieved without additive at a low temperature of 1700 \\x0e C through SPS, where the rapid densification initiates at 1189 \\x0e C. The HV and KIC of ZrC-SiC composite is 22.8 ± 0.7 GPa and 4.3 ± 0.4 MPa m1/2, respectively, which is comparable with the samples obtained under higher temperature and pressure. Furthermore, the sintered specimen shows a fine microstructure with similar embedded morphology (showing hierarchical grain size distribution), which means that the microstructure of the bulk can be predesigned or adjusted by starting powders to some extent. Researches on the characteristics of ZrCSiC ceramic with different compositions are also underway.  Fig. 14. (a) SEM image (in BSE mode) and EDS analysis of fracture surface of ZS12 ceramic sintered at 1700 \\x0e C. (b) SEM image and grain size distribution of polished surface. The orange arrows and blue arrows (the inset) point out the ZrC and SiC nano particles embedded in the domains, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)  microstructure without any pores. The micron-sized ZrC (light particle) distributes and grows along the SiC grain (dark particle) boundaries, while some tiny ZrC and SiC crystallites (about 100 nm)  \\x0c', '326  W. Xu et al.  /  Journal of Alloys and Compounds 791 (2019) 316e327  Table 2 Preparation conditions and mechanical properties of the ZrCeSiC ceramic by SPS in literature and present work.  Material  Reaction system  Sintering parameters (\\x0e C/MPa)  Relative density (%)  Vickers hardness (GPa)  Fracture toughness (MPa$m1/2)  Ref.  ZrCeSiC (~1:0.3) ZrCeSiC (~1:0.29) ZrCeSiC (~1:0.64) ZrCeSiC (~1:1.1) ZrCeSiC (~1:1.7) ZrC-SiC (not defined) ZrC-SiC (1:2) ZrC-SiC (1:2)  ZrC þ Si þ graphite ZrC þ polycarbo-methylsilane ZrC þ polycarbo-methylsilane ZrC þ polycarbo-methylsilane ZrC þ polycarbo-methylsilane ZrC þ SiC þ LiYO2 ZSi2þC ZrC þ SiC  1800/40 1950/50 1950/50 1950/50 1950/50 1600/35 1750/40 1700/45  96.1 100 99 93 94 / >99 >99  18.8 ± 1.2 27 22 / / 20.7 23.7 22.8 ± 0.7  4.0 ± 0.3 / / 3.2 ± 0.2 / 5.07 2.87 4.3 ± 0.4  [29] [9] [9] [30] [9] [28] [27] this work  Acknowledgements  The authors are grateful for the financial support from CAS Priority Research program (XDA21010204, XDB20010300), National Natural Science Foundation of China (201501178) and Natural Science Foundation of Fujian Province (2017H0048).  Appendix A.  Supplementary data  Supplementary data to this article can be https://doi.org/10.1016/j.jallcom.2019.03.299.  found  online  at  References  [8]  [6]  [2]  [1] D. Sciti, S. Guicciardi, M. 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},{
  "_id": 68,
  "PDF": "Enhanced oxidation resistance of ZrB2-SiC composite through in situ reaction of gadolinium oxide in patterned surface cavities.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 34 (2014) 4157-4166  Enhanced oxidation resistance of ZrB2 /SiC composite through in situ reaction of gadolinium oxide in patterned surface cavities  Jesus Gonzalez-Julian a,∗,1 , Omar Cedillos-Barraza b , Sven Döring c , Stefan Nolte c , Olivier Guillon a,1 , William E. Lee b  a Otto Schott Institute of Materials Research, Friedrich Schiller University of Jena, D-07743 Jena, Germany b Centre for Advanced Structural Ceramics (CASC), Department of Materials, Imperial College, London SW7 2AZ, UK c Institute of Applied Physics, Friedrich Schiller University of Jena, D-07743 Jena, Germany  Received 17 March 2014; received in revised form 23 June 2014; accepted 10 July 2014  Available online 28 July 2014  Abstract     Micro-cavities on the surface of dense ZrB2 /20 vol.% SiC composites, machined by ultra-fast laser ablation, were filled with Gd2O3 nanopowder and oxidized in static air at 1600 C. Optimized rectangular pattern of cavities, 10  \\u242em diameter and deep, 20  \\u242em apart conferred improved oxidation resistance compared  to  the untreated ZrB2 /20 vol.% SiC due  to  the  formation of glasses of higher viscosity with  lower oxygen diffusivities. Reduction of  the oxidized depth was revealed by a significant decrease of 10  \\u242em (60%)  in  the extent of  the protective  layer. The filled-cavity strategy leads to better protection against oxygen diffusivity into the composite without altering the bulk properties. © 2014 Elsevier Ltd. All rights reserved.  Keywords: Zirconium diboride; Machining; Oxidation resistance; Microstructure  1.   Introduction     Zirconium diboride (ZrB2 ) is one of the most promising materials of the ultra-high temperature ceramics (UHTCs) - borides, carbides and nitrides of groups IV and V elements  in  the Periodic Table - due  to  its unique combination of properties, such as high melting point  (3246 C), hardness  (23 GPa) and  ther−1 K −1 ),  mal  conductivity  (58.2 Wm excellent  thermal  shock (10  resistance,  low electrical  resistivity   cm) and chemical inertness against molten metals or non-basic slags.1,2 These properties have made ZrB2 an  excellent  candidate  for wing leading edges and nose  tips of  the next generation of hypersonic scramjets and aerospace vehicles where sharp profiles are required  to  reduce  the aerodynamic drag.3,4 Nevertheless,  the inherent brittleness and  low oxidation  resistance at  temperatures above 1200 C of  the monolithic material have excluded  \\u242e\\x01     ∗  Corresponding author. Tel.: +49 2461614217; fax: +49 2461615700.  E-mail address: j.gonzalez@fz-juelich.de (J. Gonzalez-Julian). 1 Present   and Climate Research   for Energy   Institute   address:   (IEK-1),  Forschungszentrum Jülich GmbH, Jülich D-52425, Germany.  http://dx.doi.org/10.1016/j.jeurceramsoc.2014.07.015  0955-2219/© 2014 Elsevier Ltd. All rights reserved.  monolithic ZrB2 from  these applications. To overcome  these problems, 20 - 30 vol.% of SiC particles have been  incorporated  into pure ZrB2 to enhance  the mechanical properties at room and high  temperatures, so  that flexural strength at room temperature exceeds 1000 MPa and fracture toughness achieves values of 5.3 MPa m1/2 .5 In addition,  the critical oxygen diffusivity  into  the pure material  is  reduced by  the addition of  the SiC second phase due to the in situ formation at high temperature of a liquid borosilicate glass, which collects at the surface, resulting from  the oxidation of ZrB2 and SiC. This glass  layer acts as an effective protective layer against oxygen penetration 1500 up  to  C.6-9 Nevertheless,  there are  still a  few challenges before implementing ZrB2 /SiC composites in hypersonic flights: (i) to develop composites with enhanced oxidation resistance at higher  temperatures, and  (ii)  to eliminate or at  least reduce the amount of liquid protective layer to prevent its being blown off in the turbulent air stream during flight. A common approach for enhancing the oxidation resistance of ZrB2 is the addition of metal oxides, silicides, borides and/or carbides within the bulk ZrB2 /SiC.10 At high temperatures metal cations are incorporated into the borosilicate glass, inducing liquid  immiscibility and forming phase separated glasses of high           \\x0c', '4158   J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166                 \\u242em   viscosity.11-15 This  immiscibility  reduces  the permeability  to oxygen by suppressing evaporation of B2O3 from the glass. As a result, 10 mol% of CrB2 , NbB2 , VB2 , TaB2 or TiB2 were added to ZrB2 /SiC composites  to enhance oxidation resistance, with successful results at 1400 C for all compounds, TaB2 being the most effective additive.16 However, Hu et al.17 reported a detrimental effect in the oxygen diffusivity at 1800 C for ZrB2 /SiC composites containing 10 vol.% of TaB2 , TiB2 , LaB6 or AlN. Incorporation of 20 vol.% TaSi2 improved  the composites oxidation resistance above 1600 C due  to phase separation  in  the amorphous surface layer, but negative effects were observed at C due to the melting of either Ta2O5 or Ta2O5 ·6ZrO2 .18 1927 Despite  the  large number of publications  in  this field,  the role of  the different metal cations  is not well understood, although a  trend has been  identified. Addition of metal cations reduces the oxygen diffusion  at  intermediate  temperatures by  liquid immiscibility whereas  the excessive  liquid formation at higher temperatures generates a negative effect due  to melting of  the solid skeleton ZrO2 -based compounds. Another interesting and novel approach is the in situ formation of an outer dense protective layer at high temperature.19,20 Some  of  the  present  authors  have  previously  incorporated 10 wt.% of rare earth additives (LaB6 , La2O3 and Gd2O3 ) into ZrB2 /20 vol.% SiC and oxidized  it at 1600 C.19 As a  result, 250  a  thick protective dense  layer of  solid  rare  earth -  zirconate  and ZrO2 was obtained  at  the  surface. A  similar composition (ZrB2 /20 vol.% SiC/10 vol.% LaB6 ) was tested at 2400 C by  acetylene  torch  and  the  in  situ  formation of LaZr2O7 and ZrO2 significantly  improved  545  the oxidation resistance due to the formation of a compact  \\u242em thick layer.20 Rare earth zirconates have been considered  for  thermal barrier coatings due  to  their high melting points  (>2300 C) and low thermal conductivity (1.5 W/mK at 1200 C).21-23 Despite these excellent results concerning oxidation resistance, the formation of a  thick continuous dense  layer based on compounds with  low  thermal diffusion modifies  some critical properties such as  the  thermal stress field  induced by  thermal diffusion under extreme conditions. For example, in a rocket motor nozzle  the  temperature at  the  inside  surface  reaches 2000 C  in less  than 1 second, whereas  the outer surface  is still at  room temperature, generating high  thermal stresses  that could  fracture  the  component.6 Therefore,  a  fresh perspective on  the next generation of components  in supersonic/aerospace flights is  required, where  the material must present a  thin oxidation protective outer  layer  in  the absence of or with only  low  liquid content at high  temperature but without altering  the bulk material  to maintain  the  required properties provided by  the ZrB2 /SiC. The  aim of  this work  is  the  in  situ  formation of  a  thin protective  layer  against oxygen diffusivity with  low  content of  liquid  at  high  temperature. By  only modifying  the  surface,  the bulk ZrB2 /SiC  remains unaltered. For  this purpose, micro-cavities have been produced on  the surface of composite specimens by ultra-fast laser ablation and filled with Gd2O3 nanopowder. Oxidation of  the specimens at 1600 C  in air was carried out  to  reveal  the  effect of Gd2O3 on  the oxidation resistance.                 2. Experimental procedure            1   ±  ZrB2 powder  (>99.0%, particle size - 400  \\u242em mesh; ESK Ceramics GmbH &Co., Kempten, Germany)  and SiC powder (␣-SiC, >99.0%, d50  \\u242em; ESK Ceramics GmbH &Co., Kempten, Germany) were used as starting materials. The asreceived ZrB2 powder was previously dry-milled in a planetary ball mill (PM100, Retsch GmbH; Haan, Germany) for 1 h using WC balls  to homogenize and reduce  the starting particle size. The final average particle size was 9.3   1.5  \\u242em, measured by laser diffraction (Malvern Nanosizer; United Kingdom). Appropriate amounts of  the  starting powders were mixed  together by mechanical  stirring  in ethanol  for 6 h  to obtain  the final composition ZrB2 /20 vol.% SiC. Afterwards, the obtained suspension was dried in a rotary evaporator and in an oven at 110 C overnight. Finally, the dried powder was sieved through 100  \\u242em mesh. Densification was carried out using  the field-assisted  sintering  technique  (FAST, Type HP D 5, FCT Systeme GmbH; Rauenstein, Germany),  also known  as Spark Plasma Sintering. 7.5 g of  the ZrB2 /SiC composite powders were  loaded  in a standard graphite tool of 20 mm inner diameter. Previously, a graphite foil was placed between  the die/punches and powder to ensure electrical and  thermal contacts. In addition, a carbon felt surrounded  the graphite die  to reduce heat  loss. A heating rate of 100 C/min to the maximum temperature (2000 C), was followed by a holding time of 10 minutes. Vacuum atmosphere of 4 Pa and uniaxial pressure of 50 MPa were maintained during the whole experiment. Temperature was measured with a vertical pyrometer focused on  the bottom of  the upper punch at a distance of 5 mm from the sample. Disc specimens of 20 mm of diameter and 5 mm  thickness were obtained. Apparent density was measured by Archimedes method and  theoretical density was calculated using the rule of mixtures. Three different patterns of cavities were produced on the surface of the dense composites by ultra-short pulse laser ablation. Ultra-short laser pulse was chosen due to the superior precision for  the processing of micro-sized  features with high  level of reproducibility.24 The pulse duration of a few picoseconds limits the ablation by energy transfer just to the size of the irradiated volume, minimizing the thermal and mechanical damage of the surrounding material.  In addition, a  large number of cavities (millions) can be formed  in a few minutes, so  that  this  technique  is  ideal for  the pattern production. Here, we used a  laser system with pulse duration of 8 ps at a wavelength of 1030 nm (TruMicro 5050, Trumpf; Ditzingen, Germany) which delivers maximum pulse energy of 125  \\u242eJ at pulse repetition rates up to \\u242eJ were focused on 400 kHz. Laser pulses with energy of ca. 19  \\u242em, correthe sample surface, with a focal spot size of ca. 15  sponding to an applied fluence of approx. 21 J/cm2 . To reach the required depth, 200 pulses were applied for each cavity at a pulse repetition rate of 10 kHz which results  in a process  time of ca. 20 ms per cavity. Depths and diameters of the cavities were kept constant for the different rectangular patterns, as only the inter cavity separation was modified between 10, 20 and 50  \\u242em by a high-precision multi-axis positioning system (ANT 130-XY, Aerotech; Pittsburgh, USA). From now,  the specimens will be  \\x0c', 'J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166   4159  Fig. 1.   (a) Secondary electron and (b) back-scattered electronimages of polished surface of ZrB2 /SiC composite. Black arrows in (b) indicate microcracks.  ±  is homogeneously distributed within  the ZrB2 matrix  (lighter grey phase), although small agglomerations of SiC are detected (Fig. 1). These clusters measured up  to 30  \\u242em, which  is above the microcracking  threshold of 11.5  \\u242em for ZrB2 /SiC reported by Watts et al.25All the thermalmicrocracks, created during the cooling, originate from SiC clusters and either finish in the ZrB2 matrix or  link  together connecting  two clusters. The presence of microcracks negatively affects mechanical properties such as elastic modulus, flexural strength and fracture toughness at room and high temperature.26 The grain size of ZrB2 is 11.2  \\u242em,  2.3  showing a limited grain growth due to the pinning effect of SiC particles. In addition, characterization by XRD (Fig. 2) revealed only two phases (ZrB2 and  ␣-SiC) suggesting formation of secondary phases during sintering is not large. Once dense ZrB2 /20 vol.% SiC was obtained,  three different patterns of micro-cavities were made (Fig. 3) with constant cavity size but varying distance between  them. Cavities  in all samples were well aligned. Distances among cavities were 10 (Fig. 3a), 20 (Fig. 3b) and 50  \\u242em (Fig. 3c) for each pattern and the cavity shape was unaffected by  the presence of SiC phase. Identical cavities were obtained for the three patterns as can be seen in Fig. 3d, e and f for ZS10, ZS20 and ZS50, respectively. Cavities of 10.0   1.5  \\u242em diameter are surrounded by affected ((cid:5)1 J/cm2 ) during material  likely melted by  the high fluence   ±  labelled as a function of the distance between cavities, i.e. ZS10, ZS20 and ZS50. Pattern dimensions were measured with a field emission scanning electron microscope (Auriga FE-SEM, Zeiss; Oberkochen, Germany), whereas depth of cavities was evaluated after ablation by  focused  ion beam  (FIB) attached  to  the FE-SEM. In addition, chemical analysis of cavities was carried out using an energy dispersive spectroscopy (EDS) unit (Oxford Instruments; Oxfordshire, UK) included in the FE-SEM. A dry process was used to fill the cavities with Gd2O3 powder (99.8%, particle size <50 nm (XRD); Sigma Aldrich). The sample with the empty cavities was pressed into a bed of nanopowder followed by a sweep process on the surface of the ZrB2 /SiC with a plastic rod. This process was repeated 10 times per sample to ensure cavity filling, which was confirmed through observation in  the FE-SEM and analyzed chemically by EDS. In addition, polished ZrB2 /SiC composites were coated with Gd2O3 by a dip-coating process as a reference material, labelled as ZSDIP. 0.5 M  suspension of Gd2O3 was homogenized  in ethanol by mechanical  stirring  for 1 h. After a dipping  time of 30 s,  the specimen was pulled continuously  through  the suspension at a constant speed of 10 cm/min. Finally, samples were dried for 1 h at 90 C. Samples and a blank of unaltered ZrB2 /SiC  (hereafter  referenced as ZS) were placed  in an alumina crucible and held at 1600 C  for 1 h  in static air. Phase analysis of sintered and oxidized samples was carried out by X-ray diffraction  (XRD, PW7100 Philips; Eindhoven, The Netherlands) using CuK␣ radiation. Samples were cut perpendicular  to  the oxidized surface  in  ethanol media  to  avoid  further oxidation,  and  cross sections were polished  to a 1/4  \\u242em finish using diamond abrasives. Thickness of the different layers through the cross section was measured using FE-SEM micrographs, and  their chemical composition using EDS. Chemical analyses of the oxidized surfaces were carried out using  time of flight  secondary  ion mass spectrometry  (ToF-SIMS, TOF.SIMS5 ,  IONTOF; Münster, Germany). Samples were held in the analysis chamber under ultra-high vacuum and bombarded their surfaces with Caesium ions.        3. Results and discussion  Dense specimens of ZrB2 /20 vol.% SiC  (99.3%  theoretical density) were obtained after sintering by FAST. SiC (dark phase)  Fig. 2. XRD of starting ZrB2 and SiC powders and sintered ZrB2 /SiC composite.  \\x0c', '4160   J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166  Fig. 3. FESEM micrographs of ZS10 (a and d), ZS 20 (b and e) and ZS50 (c and (e) patterns.  the laser process. In the vicinity of the cavities an outer ablation zone was observed, especially  in  the case of ZS20. This area is probably produced by  the distortion of  the beam profile and shows typical ripple formation at low fluence (<1 J/cm2 ). Unaltered material can be easily distinguished. Cavity depth was measured using a FIB (Fig. 4a) and  the compounds created by the laser ablation were characterized using EDS (Fig. 4b). The cavities were 10  \\u242em deep in all samples and no degradation such as microcracks was observed in the surrounding material. Chemical analysis of  the molten material and  the outer ablation zone  (points  I and  II)  reveals zirconium  (Zr), boron  (B), oxygen (O), silicon (Si) and carbon (C) suggesting formation of  ±   1   a borosilicate glass from the degradation of ZrB2 and SiC during the ablation process. In contrast,  the untreated material (point III)  reveals  just Zr and B  (from ZrB2 ) discarding  the surface oxidation and confirming  the precision and  innocuousness of the ultrafast laser technique for machining patterns on ZrB2 /SiC composites. The success of  the dry process  for filling  the cavities with Gd2O3 nanopowder was confirmed using FE-SEM and EDS (Fig. 5). All the cavities are completely filled with the nanopowder with a  small amount adhering on  the  surface around  the hole (Fig. 5a-c). EDS analysis confirms that cavities were filled with Gd2O3 (Fig. 5d). After  the filling process, five different  Fig. 4.   (a) FIB image of cross section and (b) FE-SEM micrograph of a cavity in ZS20. EDS were taken on regions marked (I-III).  \\x0c', 'J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166   4161  Fig. 5. FE-SEM micrographs of Gd-filled ZS20 (a-c) and EDS on region I (d).     samples were oxidized for 1 h  in static air at 1600 C:  three of them were ZS10, ZS20 and ZS50 filled with Gd2O3 , and two reference specimens, a blank of unaltered ZrB2 /SiC composite and another one of ZrB2 /SiC coated with Gd2O3 that was processed by dip coating  (ZS and ZSDIP,  respectively). Phase analysis of all  the sample surfaces after  the oxidation process  is shown in Fig. 6. XRD detected  the  same phase composition  for all the specimens, predominantly m-ZrO2 and trace of t-ZrO2 . No presence of Gd-crystalline compounds was observed by XRD, discarding  the  in situ  formation of gadolinium zirconate during  the oxidation process  for ZS10, ZS20, ZS50 and ZSDIP. Figs. 7 and 8 show representative surface and cross section pictures, respectively, of all the oxidized samples. ZS presented the characteristic oxidized surface (Fig. 7a and b), spherical particles  Fig. 6. XRD of oxidized ZS, ZS10, ZS20, ZS50 and ZSDIP.  of ZrO2 embedded into the borosilicate glass. These ZrO2 particles were precipitated  from  the borosilicate glass and  their concent at the top of the surface covered a wide area. Incorporation of Gd2O3 into the cavities (Fig. 7c-h) and by dip coating (Fig. 7i and j) clearly modified the surface microstructure of the oxidized specimens. ZS10, which presented the highest amount of Gd2O3 , was completely covered by the borosilicate glass and the content of ZrO2 particles was scarce (Fig. 7c). Interestingly, phase separation phenomenon was observed, although its presence was marginal (Fig. 7d). Reduction of  the Gd2O3 content increased the precipitation of ZrO2 particles, ZS20 presented a major concentration  than ZS10 (Fig. 7e and f), whereas ZS50 surface microstructure was similar to ZS (Fig. 7g and h). Finally, ZSDIP oxidized as ZS10, surface was completely covered by the borosilicate glass and only some precipitated ZrO2 particles were observed (Fig. 7i and  j). Once  the oxidized surfaces were analyzed, cross sections of all  the specimens were evaluated to ascertain the oxygen diffusivity through the specimens. Microstructure of ZS (Fig. 8a) presents the characteristic layers as previously described.9-12,14-19 A continuous and homogeneous 25  \\u242em thick outer layer of borosilicate glass is on top of an intermediate 75  \\u242em thick porous layer containing ZrO2 particles and pockets of borosilicate glass. Finally,  the unaltered ZrB2 /SiC can be easily observed. Cross sections of ZS10, ZS20 and ZS50 are shown in Fig. 8b, c and d, respectively, and present the same layers but with different thicknesses. Interestingly, the external  layer  is a non-uniform mixture of a glass phase and large numbers of particles with sizes between 0.7 and 4.0  \\u242em, especially for ZS10 and ZS20, which are difficult to distinguish from the intermediate layer. For these specimens the thickness of glass in the outer layer (Fig. 8b and c) is considerably lower than for ZS (Fig. 8a), some outer areas were even glass free. Estimates of  the  thickness of  this  top  layer were ca. 20   3, 15   3, and 30  \\u242em for ZS10, ZS20 and ZS50, respectively. Regarding  ±  ±  ±   2   \\x0c', '4162   J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166  Fig. 7. Surface micrographs of specimens oxidized at 1600  in (d) indicate phase separation (PS).     C for 1 h (a and b) ZS, (c and d) ZS10, (e and f) ZS20, (g and h) ZS50 and (i and j) ZSDIP. White arrows  the  intermediate  layer, based mainly on ZrO2 particles, pockets of glass and porosity,  the difference  in scale  is significant. Porosity  is created due  to oxidation and further evaporation of the glassy phase. ZS10 and ZS50 present thicknesses of 130 and 160  \\u242em,  respectively, while ZS20  just exhibits 75  \\u242em. At  the bottom of  the micrographs,  the unaffected ZrB2 /SiC region  is easily detected. Finally, cross section of ZSDIP (Fig. 8e) shows practically  20  the same structure as  the other specimens, an outer layer of  \\u242em based on glass and 100  \\u242em thick intermediate porous  layer of ZrO2 and glass between  the external  layer and the underlying ZrB2 /SiC composite. Distance between  the filled cavities plays a critical  role  in the composition of  the outer  layer and  the oxygen diffusivity according to the results obtained after oxidation of ZS10, ZS20 and ZS50 (Figs. 7 and 8). The response to oxidation depends on the starting composition and  its evolution during  the reactions  at high  temperature. The main  reactions expected during  oxidation at 1600 C are shown below6,8,10,19,20 :     the  ZrB2 (s)   +  SiC(s)   +   5/2O2 (g)   3/2O2 →  →   ZrO2 (s)   +   B2O3 (l)    SiO2 (l)   +   CO (g)  B2O3 (l)   B2O3 (l)   →   B2O3 (g)   +   SiO2 (l)   →   borosilicateglass   (1)  (2)  (3)  (4)  (5)  (6)  When Gd2O3 is incorporated into the cavities, the following reactions can also occur at high temperatures:  Gd2O3 (s)  Gd2O3 +  +   ZrO2 (s)   →   Gd2Zr2O7 (s)    SiO2 (s)   +   B2O3 (l)   →   Gd - borosilicateglass   \\x0c', 'J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166   4163  Fig. 8. Cross section micrographs of specimens oxidized at 1600     C for 1 h (a) ZS, (b) ZS10, (c) ZS20, (d) ZS50 and (e) ZSDIP.     During the heating, ZrB2 and SiC oxidize through reactions (1) and  (2)  forming ZrO2 , boria  (B2O3 ), SiO2 and CO. As a result, a solid oxide, ZrO2 , is surrounded by a liquid borosilicate glass, formed by reaction (4). This liquid flows from oxidation sites towards the surface, spreading the whole surface. Borosilicate glass  is an effective protective outer  layer  to oxidation due  to  it has  lower oxygen diffusivity  than silicate glass and a higher viscosity and boiling point  than boria. However,  the amount of B2O3 is reduced from 1200 C due  to  its volatilization  (reaction  (3)). As a  result,  the exact composition of  this glassy phase  is undefined. Borosilicate protective glassy phase is clearly observed in ZS (Fig. 8a), and hinders the diffusion of oxygen  into  the bulk material. In addition, oxygen penetration was also  identified  in  the  layer beneath due  to  the presence of spherical ZrO2 particles embedded in a small volume of borosilicate glass in the bulk material, formed by the oxidation of ZrB2 , through reactions (1) and (2). Oxidation response of  the material  is modified by  the addition of Gd2O3 into  the cavities, which could  trigger  reactions (5) and (6) by the incorporation of new cations into the complex glassy phase. Fig. 9 presents a FE-SEM detail of  the external glassy phase for ZS20 and an EDS analysis. Spherical particles were embedded into the glassy phase (Fig. 9a). Elemental maps by EDS reveal the distribution of Zr, Si and Gd in the outer layer (Fig. 9b, c and d, respectively). Zr  is observed  in  the spherical  particles, while Si is mainly detected in the glassy phase. Intensity of Gd is low and ambiguous, and it is randomly detected in both phases (Fig. 9e and f). Therefore, the ambiguous location of Gd atoms by EDS required the precise analysis of the oxidized surface by ToF-SIMS (Fig. 10). Elemental maps of silicon (Si+), boron (B+), gadolinium (Gd+) ions and their overlay image are illustrated in Fig. 10a, b, c and d, respectively. Gadolinium was homogeneously distributed into the glass phase, which is composed by boron, silicon and gadolinium. As a result, gadolinium ions were incorporated to the borosilicate glass. In addition, the overlay image (Fig. 10d) of the ions reveals that the glass was not completely homogeneous, presenting boron - rich and - poor areas. As  it was  previously  revealed  by XRD,  and  confirmed by FE-SEM  and ToF-SIMS,  gadolinium  zirconate was  not in situ  formed during  the oxidation process, discarding  reaction (5) for our oxidation conditions. Nevertheless, gadolinium zirconate  formed by  reaction  (5)  at 1600 C  in  air has previously been  reported.19 Conditions  for  in  situ  formation of zirconates at high  temperature are not perfectly understood as report.19,20,27 Success or  the  few published papers  failure of in situ zirconate formation depends on the molar ratio between zirconia and rare-earth oxide (reaction (5)), kinetic and thermodynamic factors and starting powders,  i.e. rare earth oxides or borides.27     \\x0c', '4164   J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166  Fig. 9. Outer layer of ZS20 oxidized at 1600  a spheroidal particle (f).     C for 1 h. FE-SEM micrograph (a), EDS maps for Zr (b), Si (c), and Gd (d), and EDS spectra from glass phase (e) and  The  role  of  a metal  cation  in  the  glassy  phase  is well understood due  to  the  larger number of publications on  the addition of metal cations  to  the borosilicate glass present  in oxidized ZrB2 /SiC composites.11-18 These foreign atoms cause liquid immiscibility and phase separation in glasses, leading to higher viscosity and lower oxygen diffusivities. This effect was the addition of different cations such as, Cr3+ , observed with  Ta5+ , Ti4+ , V4+ and Nb5+ , and for  temperatures from 1400  to 1627 C.18 On  the other hand, a detrimental effect on  the oxidation penetration by  the addition of  rare-earth oxides occurs by facilitating  the  inward diffusion of oxygen by  the  increased     concentration of oxygen vacancies  in  the zirconia particles.27 Therefore, meticulous analyses of  the oxidized  surfaces and cross sections of all the specimen are required to reveal the role of the gadolinium in the oxidation rate due to it presents beneficial and detrimental phenomena  in  the oxidation response of the ZrB2 /SiC composite. In our case, the role of the immiscibility observed for ZS10 was considered of marginal  importance due  to  its  low concentration. Same marginal effect was reported by Monteverde et al. in  the  inward oxygen diffusion rate  through  the oxide scale.27 The beneficial effect of the incorporation of Gd2O3 was related  \\x0c', 'J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166   4165  Fig. 10. ToF-SIMS maps of (a) silicon, (b) boron and (c) gadolinium cations, and (d) overlay of the elemental ions maps.  with the increment of the glass viscosity and lower oxygen diffusivities. According  to  these assumptions, ZS10 must present the thinner intermediate oxide scale in the cross section pictures (Fig. 8), followed by ZS20 and ZS50. However, the thinner oxide scale corresponded to ZS20. There are two factors that have to be considered for an exhaustive analysis, (i) the oxygen vacancies in the zirconia particles by the addition of gadolinium ions and (ii) the oxidant protection by the precipitated ZrO2 particles. The most critical effect is the formation of oxygen vacancies in the zirconia particles, leading a higher oxidant rate through the particles. This response can be easily observed comparing the cross sections of ZS and ZS50 (Fig. 8). Both oxidized microstructures of  the outer  layer were similar, high concentration of precipitated ZrO2 particles embedded  into  the glass phase  that had a thickness around 20  \\u242em. Nonetheless,  the  intermediate oxide scale is larger for ZS50 and just could be related with the effect of  the gadolinium  in  the glass phase and  its detrimental phenomenon by formation of oxygen vacancies into ZrO2 particles. Another important effect is the oxidation protection supplied by the precipitated ZrO2 particles, although  their quantification  is rather difficult. As a  result,  the effect of  the gadolinium  ions in the oxidation behaviour of ZS10, ZS20, ZS50 and ZSDIP is described below. The highest concentration of gadolinium  in  (ZS10)  implied  a  high  viscous  glass with   the outer  layer lower  oxygen  diffusivities  that protected  the bulk material against  the oxygen penetration and  reduced  the amount of  liquid content of the external  layer. However,  the high gadolinium content hindered  the ZrO2 precipitation at  the  top of  the glass phase and generated high concentration of oxygen vacancies  in  the ZrO2 particles, leading the inward oxygen diffusion through the material. As a result, the content of glass phase is reduced in the outer layer but  the  inward oxygen diffusion  is  increased  in comparison with  the unaltered ZrB2 /SiC composite. Same behaviour is assumed  for ZSDIP. Reduction of  the gadolinium content in ZS20 specimen maintained  the high viscosity and  the  lower oxygen diffusivity, keeping  the  low content of glass. However, the  lower amount of gadolinium allowed  the precipitation of ZrO2 particles and reduced oxygen vacancies concentration  in the ZrO2 particles. Consequently, the outer and the intermediate layers were reduced in comparison with ZS10. Finally, the low content of gadolinium ions in ZS50 implied practically the same outer  layer  in  terms of glass viscosity, ZrO2 precipitation and thickness than ZS. However, the detrimental generation of oxygen vacancies in ZrO2 particles generated higher oxidation rates than ZS. According to the experimental results, ZS20 is selected as optimized filled pattern due to the high viscosity of the glass and  the reduction of  the outer  layer with  the minimum oxygen inward diffusion. Difficulties in estimating packing in the cavities hinder a rigorous comparison between ZS20 and ZSDIP due  \\x0c', '4166   J. Gonzalez-Julian et al. / Journal of the European Ceramic Society 34 (2014) 4157-4166  to the difference in the initial content of Gd2O3 . Despite the better results from ZS20, further research of ZrB2 /SiC composites coated with different amounts of Gd2O3will be appropriate. For ZS20 the presence of gadolinium in the borosilicate glass reduces  the  thickness of  the outer protective  layer  in a 60%  in \\u242em) and ZS20 (15  \\u242em). The  comparison between ZS (25  total oxide  thickness, outer and  intermediate  layers,  is  reduced by 10% with  the  incorporation of  the gadolinium  in comparison with the unaltered ZrB2 /SiC. Interestingly, the most remarkable observation is the reduction of glassy phase content in the outer layer. The processing method developed here and based on the addition of powder in cavities for in situ formation of a protective layer during  the  thermal process suggests an alternative  route for development of materials for high temperature applications. Further research  is required  to  test  the resulting composites at high temperature and in air streams.  4. Conclusion     Dense  ZrB2 containing  20 vol.%  SiC  composites were obtained by FAST and machined by an ultra-fast  laser ablation process to create three different patterns of cavities on their surfaces. Micro-cavities, 10  \\u242em diameter and deep, were filled with Gd2O3 nanopowder and oxidized  in static air at 1600 C. \\u242em Composites with  an  array  of  cavities  separated  by  20  showed better oxidation resistance  than unmodified ZrB2 /SiC. Gadolinium was  incorporated during  the  thermal process  into the borosilicate glass, entailing  liquid  immiscibility and phase separated glasses of higher viscosity and lower oxygen diffusivities. However, detrimental effect in the inward oxygen diffusion was observed with  the  incorporation of gadolinium  ions  into the glass phase due  to  the  formation of oxygen vacancies  in ZrO2 particles. A optimized control of  the gadolinium oxide content  implies a significant decrease of 60%  in  the  thickness of protective  liquid  layer  that spreads over  the ZrB2 /SiC. This method opens the possibility for using ZrB2 /SiC composites in real applications under  stream air  such as  in hypersonic and aerospace flights.  Acknowledgments  This work is financially supported by the European Ceramic Society (ECERs) under  the “Frontiers of research” project for young  researchers. Raw ZrB2 and SiC powders were donated by ESK Ceramics GmbH & Co (Germany).  References  1. Zimmermann JW, Hilmas GE, Fahrenholtz WG. Thermophysical properties of ZrB2 and ZrB2 -SiC ceramics. J Am Ceram Soc 2008;91:1405-11. 2. Guo S. Densification of ZrB2 -based composites and  their mechanical and physical properties: a review. J Eur Ceram Soc 2009;29:995-1011.  3. Wie D, Drewry D, King D, Hudson C. The hypersonic environment: required operating conditions and design challenges. 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Oxidation of ZrB2 ceramics modified with SiC and group IV-VI transition metal diborides. Electrochem Soc Proc 2001;12:144-58.  17. Hu P, Zhang X, Ha J, Luo X, Du S. Effect of various additives on the oxidation     C. J Am  behavior of ZrB2 based ultra-high-temperature ceramics at 1800 Ceram Soc 2010;93:345-9.  18. Opila E, Levine S, Lorincz J. Oxidation of ZrB2 and HfB2 based ultra-high temperature ceramics: effect of Ta additions. J Mater Sci 2004;39:5969-77.  19.   Jayaseelan DD, Zapata-Solvas E, Brown P, Lee WE. In situ formation of  oxidation  resistant  refractory coatings on SiC  reinforced ZrB2 ultra high temperature ceramics. J Am Ceram Soc 2012;95:1247-54.  20. Zhang X, Hu P, Han J, Xu L, Meng S. The addition of   lanthanum hexa boride to zirconium diboride for improved oxidation resistance. Scr Mater 2007;57:1036-9.  21. Belmonte M. Advanced ceramic materials for high temperature applications. Adv Eng Mater 2006;8:693-703.  22. Wu J, Wei X, Padture N, Klemens P, Gell M, Garcia E, et al. Low   ther mal conductivity rare earth zirconates for potential thermal barrier coating applications. J Am Ceram Soc 2002;85:3031-5.  23. Zhao H, Levi C, Wadley H. Vapor deposited samarium zirconate  barrier coatings. Surf Coat Technol 2009;203:3157-67.  thermal  24. Samant A, Dahotre N. Laser machining of structural ceramics - a review. J Eur Ceram Soc 2009;29:969-93.  25. Watts   J, Hilmas GE, Fahrenholtz WG. Mechanical   characterization of  ZrB2 -SiC composites with varying SiC particle  2011;94:4410-8.  sizes. J Am Ceram Soc  26. Neuman EW, Hilmas GE, Fahrenholtz WG. Mechanical behavior of zirco nium diboride-silicon carbide ceramics at elevated temperature in air. J Eur Ceram Soc 2013;33:2889-99.  27. Monteverde   F, Alfano D,   arc-jet   convectively   perature  ceramics  2013;75:443-53.  in   R.   Effects   of   Savino   LaB6 heated SiC-containing ZrB2 -based  ultra-high  high  enthalpy  supersonic   airflows. Corrosion   addition   on  tem Sci  \\x0c']"
},{
  "_id": 69,
  "PDF": "Evaluation of oxidation behaviors of HfC-SiC ultra-high temperature ceramics at above 2500 °C via oxyacetylene torch.pdf",
  "Text": "['Ceramics International 44 (2018) 8505-8513  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Evaluation of oxidation behaviors of HfC-SiC ultra-high temperature ceramics at above 2500 °C via oxyacetylene torch  T  Young-Hoon Seonga,b,1, Changyeon Baeka,1, Joo-Hyung Kima, Jung Hoon Konga, Dong Seok Kima,c, Sea-Hoon Leed, Do Kyung Kima,⁎  a Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology (KAIST), Yuseong-gu, Daejeon 34141, Republic of Korea b Korea Institute of Energy Research (KIER), Yuseong-gu, Daejeon 34129, Republic of Korea c Korea Atomic Energy Research Institute (KAERI), Yuseong-Gu, Daejeon 34057, Republic of Korea d Korea Institute of Materials Science (KIMS), Changwon, Gyeongnam 51508, Republic of Korea  A R T I C L E  I N F O  A B S T R A C T  Keywords: HfC-SiC UHTCs Oxidation behavior Oxyacetylene torch  1.  Introduction  Development of materials highly resistant to harsh conditions (including ultra-high temperature) has been a crucial issue for the development of next-generation hyper velocity vehicles. HfC-SiC ceramics have been attractive as suitable candidates for these applications owing to their extraordinarily high melting temperatures. Here, HfC-SiC ceramics with ultra-fine grains, prepared via reactive spark plasma sintering, were tested under an oxyacetylene torch flame for relatively long exposure times (5-30 min). The oxidation behaviors of the samples were carefully characterized by microstructural analysis. Although porous structures were formed due to the generation and escape of gas phases, the mixture of oxidized species composed of HfO2 granules and melted SiO2 phase, acted as an excellent barrier under the condition of severe high-temperature oxidation.  Increasing interest in re-entry hypersonic vehicles and weapons has driven demand for materials that might be used for leading-edge, nosetip, and propulsion-system components. The buffer materials for these applications are exposed to extreme environments including aerodynamic heating introduced by continuous friction at the interface between the material and the atmosphere. To protect the main body of projectiles, these protecting materials must withstand the severe thermal shock and oxidation that occur at ultra-high temperatures (> 2000 °C) [1,2] Among many kinds of protecting materials that survive these severe conditions, Hf-based ceramics such as HfB2 and HfC are among the most promising candidates. HfC ceramics are also known as materials that are durable at very high temperature, along with other carbide ceramics including TaC and ZrC [3]. In addition, the oxide scales (e.g., HfO2 or HfO2-x) that form as a result of the oxidation of HfC ceramics, have a very high melting temperature, over 2500 °C [3,4]. This high melting temperature of these oxidized scales enables the underlying materials to survive even in ambient (oxidizing) atmosphere. Despite the merits of HfC ceramics, they are also known to have both a low sinterability due to their high melting temperature and gas evolution problem when it is  oxidized [5-8]. Therefore, it is important that many researchers have found that addition of SiC to form HfC-SiC ceramics maximizes their thermal resistance by enhancing their sinterability [9,10]. The added SiC can also support oxidation resistance by forming SiO2 scale at lower temperature under oxidizing conditions [11,12]. For these reasons, the HfC-SiC ceramics have been investigated as various forms of composite structure, such as the inter-layered structure of HfC-SiC with BN or graphite, the multilayer coating for C-C composites, and infiltrated HfC into a C-C matrix; to avoid fracture due to the low thermal shock resistance of HfC-SiC ceramics [13-16]. It is also very important to determine the durability of such materials before making actual vehicles because a great deal of time and expense will be needed to make the final hypersonic vehicles. Therefore, evaluation of these materials must precede actual applications. High-temperature testing using an oxyacetylene torch is one of the most frequently used methods, and has the merits of simplicity, low cost, and sufficiently high temperature (~ 3000 °C) with respect to the oxygen to acetylene gas ratio [13-32]. There have been a few reports from studies to investigate the high-temperature oxidation behavior of other UHTCs under the severe exposure to oxyacetylene flame for up to 40 min [18] Here, we discuss the oxidation behaviors of nano-grained HfC-SiC ceramics under this harsh condition using a newly designed  ⁎ Corresponding author. E-mail address: dkkim@kaist.ac.kr (D.K. Kim). 1 Equal contributors.  https://doi.org/10.1016/j.ceramint.2018.02.049 Received 9 August 2017; Received in revised form 2 February 2018; Accepted 5 February 2018  Available online 06 February 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  oxyacetylene torch testing system. The time-dependent change in the oxidation scale and ablation rate, and the corresponding microstructural aspects were carefully provided to explain the oxidation behavior of fine-grained HfC-SiC ceramics prepared via reactive spark plasma sintering.  2. Experimental section  2.1. Preparation of HfC-SiC composite ceramics  HfC nanopowder (particle size: 100 - 200 nm), HfSi2 (particle size: 1-2 µm; Alfa Aesar, MA, USA) and carbon black powders (surface area: 50-80 m2/g, purity: 99.5%; Alfa Aesar, MA, USA) were used as starting materials to prepare dense ceramics. The HfC nanopowder and the HfSi2 - C mixture were mixed with the weight ratio of 6:4 by highenergy ball-milling for 2 h in dry condition using a shaker mill (Spex D8000, Spex CertiPrep, Metuchen, NJ) with tungsten carbide balls and jars to yield the well homogenized HfC-34 vol% SiC ceramics. The jars were sealed in a glove box under N2 gas to suppress oxygen contamination. The mixtures were granulated using a 150 mesh; then loaded into a graphite mold in the glove box. The mixed powders were sintered using reactive spark plasma sintering (SPS) (Dr. Sinter 2020, Sumitomo Coal Mining Co., Tokyo, Japan) at 1900 °C with a uniaxial pressure of 40 MPa for 10 min in vacuum (~20 Pa). HfC-SiC (Φ18 × 6 mm2) ceramics were obtained by in situ reaction between HfSi2 and carbon during the sintering process.  2.2. Oxyacetylene torch testing and characterization  To investigate the oxidation resistance of the prepared sample, the oxyacetylene torch set-up, which included a temperature monitoring system, a sample holder, mass flow controllers, and a sample stage was custom-built (Fig. 1a). Most of the test procedures (except torch ignition and distance adjustment) were controlled via computer to maintain reproducible test conditions. As-prepared HfC-SiC composite ceramics were polished before the high-temperature oxidation test to minimize the effect from a rough surface. The samples were joined to graphite holders and placed perpendicular to the oxyacetylene torch tip. The distance between the torch tip and the sample surface was 40 mm in all cases to minimize the heat deterioration of the nozzle tip during the tests. After adjusting the distance, the sample stage was retracted to avoid premature heating by the torch flame. When the desired torch flame was ready (O2 to C2H2 gas ratio of 1.3: O2 = 0.216 l/s; C2H2 = 0.167 l/s), the sample stage was advanced to the test position at a speed of 30 mm/s. After the test, we put the torch flame out and kept the tested sample at the test position to let it cool naturally. The heat flux delivered to the surface of the samples was measured using a heat flux gauge (TG1000, Vatell Corp., Christiansburg, VA). The gauge was introduced into the flame, a maximum voltage could be read, and the measurement is repeated three times. The heat flux of the steady oxyacetylene torch was about 800 W/cm2. This value is similar to the heat flux of the oxyacetylene torch flame reported in the literatures [24,25]. The test conditions included the distance between the samples and the nozzle, the gas flux, and the fuel to oxygen ratio. Fig. 1b includes photographs taken during the oxidation test. The surface temperature of the sample was monitored using a 2-color pyrometer (Pyrospot DSR 10NV, DIAS Infrared GmbH, Dresden, Germany) and recorded in real is 0.7 - time in the computer. The spectral range of the pyrometer 1.1 µm. The inset of Fig. 1b shows the in situ monitoring images obtained by the 2-color pyrometer during the test. The visual changes that occurred during oxyacetylene torch testing could be observed using this in situ monitoring system. To observe the oxidation behavior, samples were prepared and tested for 5, 10, 20, and 30 min. After the high-temperature oxidation test, the samples were mounted in resin and sectioned at the center of the spot exposed to the torch flame. The cross-sections of the polished samples were  Fig. 1. (a) Scheme for  the newly designed oxyacetylene torch set-up,  (b) Photographs  taken during the actual oxidation test and in situ monitoring images (inset) captured from  the recorded video files via a 2-color pyrometer system.  characterized using an optical microscope (OM; Dino-Lite Pro AM413ZT, AnMo Electronics Corp., Taiwan) and a field-emission scanning electron microscope (FESEM; Model XL30, Philips, The Netherlands) to measure the thicknesses of the residual matrix layer and the surface oxidation layer, respectively. The crystal structure of the oxidized surfaces was analyzed by X-ray diffractometer (XRD; Rigaku D/Max-RB (12 kW), Tokyo, Japan) with CuKα radiation (λ = 1.5148 Å) operating at 40 kV and 200 mA, with step size of 0.01° at 2° min−1 in a 2θ range from 20° to 70°.  3. Results and discussion  The crystal structures of the as-sintered and as-tested samples presented in Fig. 2 clearly show the surface oxidation of the HfC-SiC ceramics. The as-sintered sample mainly had a hafnium carbide phase with a small amount of silicon carbide phase. Successful in situ reaction of carbon with the added HfSi2 occurred during SPS sintering without leaving unreacted species. The surface crystal structures of the oxidation-test samples had HfO2 phase resulting from the oxidation reaction of HfC. There were no other distinguishable phases such as a crystallized forms of SiO2 or HfSiO4 (hafnon phase) which could be caused by the solid solution of HfO2 and SiO2 [16]. The sample surface was oxidized under the oxyacetylene flame and a clearly defined spot remained on the sample surface exposed to the intense flame. Fig. 3 shows the as-prepared sample and the as-tested samples. The as-prepared sample has the relative density of 100% with average grain size of 0.48 µm for HfC and 0.39 µm for SiC grains respectively. The sample also shows excellent mechanical properties  8506  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  Fig. 2. Crystal structures of  the as-sintered HfC-SiC sample and the surface oxide layers  formed on the tested sample under the flame of an oxyacetylene torch.  owing to its well-densified microstructure with nano-size grains. The detail information about the prepared HfC-SiC ceramics can be found in the previous study made by L. Feng et al. [10] The as-tested samples were fractured into several pieces by thermal shock during rapid cooling after the test. Fig. 4 shows the surface temperature profile of the prepared HfC-SiC ceramic exposed to the oxyacetylene torch flame for 30 min. The sample heated up immediately after the sample was exposed to the flame. The surface temperature reached ~2500 ℃ in a few minutes, and gradually increased to 2800 ℃ as a function of time. The ratio of radiance received by the 2-color pyrometer is independent of target size and of the position of the target within pyrometer field of view. It means that even if the central area of the specimen shone evenly, there might be a temperature gradient. If some of the spots reached a temperature of 2800 ℃ in the field of view of a 2-color pyrometer, they would have been measured at 2800 ℃ even if there were areas that did not reach 2800 ℃. We believe that the calibration of the 2-color pyrometer is reliable because reasonable results were obtained when testing other ultra-high temperature specimens such as tungstenrhenium alloy [33]. After the sample was exposed to severe temperature change, the cracks and faint crater traces were observed in the  Fig. 4. The recorded surface temperature profile of the HfC-SiC sample during oxidation  test (heating) for 30 min.  surface layer with color change from dark grey to white. Fig. 5a shows the SEM images of the tested samples with different testing time varying from 5 to 30 min. Surface-oxidized layers formed after the test. The thickness of the uneven and crusted layers increased with testing time. Crater-like morphology appeared in the oxidized layers of the samples tested for 20 and 30 min. The dark regions within the oxidized surface layers are voids, and some of the voids were filled with resin during the specimen mounting process. It is thought that the porous oxidized layers originated from severe gas evolution that occurred under the test conditions. The evolved gas phases were accumulated and exhausted through the oxide scale leaving voids during cooling. The void areas were excluded when the thickness of the oxidized layers was measured. Loss of the oxidized species could have occurred during long exposure to the flame, considering previous studies in which the oxidation behaviors of UHTCs were investigated using an oxyacetylene torch. The change in the thickness of the residual matrix also had to be analyzed to confirm whether it went along with the trend of change in the oxidation thickness. To measure the residual matrix thickness, several SEM images were taken for a sample and used together to confirm the thickness of residual thickness. This is because the normal SEM has a small field of view. To minimize the inconvenience and to promote the accuracy of measurement, an optical microscope was used because it has a larger field of view than that of the SEM. Fig. 5b shows the images captured using the optical microscope. From these images, the oxide scale can be more easily observed owing to their definite color difference. Likewise, the residual matrix appears distinct, and this helps to measure the thickness of the matrix phase. The persistence rate (or recession rate) was calculated from the  Fig. 3. SEM image of the as-prepared HfC-SiC ceramic and the photographs of as-tested samples exposed to oxyacetylene torch flame for different exposure times.  8507  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  Fig. 5. The cross-sectional SEM images (a) and OM images (b) of  the test samples exposed for different times.  difference between the original thicknesses of the samples (t0) and the residual thickness (t1) after the test, using the simple equation (t0 − t1)/ t0 * 100. The information obtained about the oxidation behavior of the HfCSiC samples is summarized in Fig. 6. The variation of the oxidation thickness with the heating time (Fig. 6a) shows a parabolic trend. This reduction in the oxidation rate can also be confirmed from the change in residual thickness shown in the inset of Fig. 6b. The thickness of the observed oxidation layer and the recession rate after 30 min exposure, were comparable to the values reported from oxidation tests of different kinds of UHTCs under oxyacetylene torches as shown in Table 1 [13,18,28,31]. This suggests that the oxidation of HfC-SiC under the oxyacetylene torch flame follows reaction diffusion-controlled kinetics [18]. In the low magnification SEM images (Fig. 5a), the surface layer seems porous. From the higher magnification SEM images (Figs. 7 and 8), however, it can be noticed that the surface layer of tested samples is partially densely composed of HfO2 granules and SiO2 liquid phase acted as a protective barrier to prevent oxygen diffusion into the HfCSiC matrix. The parabolic trend in the oxidation behavior generally occurred when the diffusion of oxygen was slowed down due to the presence of the dense surface oxidized layer. The linear ablation rate  (Fig. 6b) calculated by dividing the thickness of the residual surface oxidation layers by the heating time could be reduced owing to the presence of this protective surface layer. Surface and cross-sectional SEM images were used to analyze the oxidation behavior of the HfC-SiC ceramics. Fig. 7 shows the surface morphology of the sample tested for 20 min. Among all the tested samples, the oxidation layer of the sample tested for 20 min was chosen because the oxidation rate declined at this time, and because the sample had features that appeared during both immediate and late periods of the test. There was a common characteristic that the surfaces of the oxidized crusts were porous in all the samples tested. Overall, the samples had rough surface morphology with two distinguishable phases. The main phase (dark) was smooth and wavy without any noticeable boundaries. The other phase (bright) had pebble-like granules a few microns wide. The EDS elemental analysis revealed that the dark area (spot 1) was composed of Si and O, which means that SiO2 phase was formed the during oxidation test. The SiO2 liquid phase was consolidated in amorphous phase due to the rapid cooling (quenching) after the test. This could be the reason that the crystallized SiO2 phases were not identified in the XRD patterns. On the other hand, the bright granules (spot 2) consisted of Hf and O, which implies that these  8508  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  As a result of observing the cross-sectional microstructure of the surface layer (Fig. 8(a), region 1), it was found that liquid SiO2 was distributed directly below the surface in most regions. The solid phase HfO2 in the form of spherical particles (Fig. 7(a)) and molten HfO2 that formed a large neck (Fig. 7(b)) were also found from the microstructures of surface. The cross-sectional SEM images (Fig. 8a) of the tested sample were carefully scrutinized to analyze changes in the aspect of the oxidation layer with respect to the depth from the surface. The dark areas shown at the top, and in the pores of the oxidation layer, are polymer resin that filled in during the mounting process. The oxidation layer was subdivided and magnified into three definite regions from the surface to the matrix (near surface, middle, and near matrix). Considering the images taken in secondary electron mode, the brighter region is composed of heavier elements. Taking into account of the contrast and the EDS analysis obtained from the surface images, the bright granules and the dark phase are HfO2 and SiO2 phases, respectively. Based on above results, it is apparent that the bright and fine HfO2 granules are distributed within SiO2 in Region 1, which is located near the surface. Some of the SiO2 phase is also consolidated and cracked in some parts of Region 1. In Region 2, which is located between the surface and the matrix, the SiO2 phase is mostly located in the upper part (near the surface) and is less evident in the bottom part (near the matrix). Instead, most of the bottom part is HfO2 with micro-pores. The slightly darker layer in Region 3 is an interlayer located between the matrix layer and delaminated surface layer. The presence of an interlayer is more noticeable in the optical image of the slightly polished surface. When the white surface was removed by polishing, the black-colored interlayer was exposed between the surface oxide and the matrix as shown in Fig. 8b. Based on the EDS results obtained from two different spots (1 and 2) in the BSE image in Fig. 8b, the interlayer consists of Hf, C, and O without Si, although a slight contrast difference is present in the image. The flux of the melted SiO2 phase from near matrix to outermost layer of the oxidized surface caused the depletion of Si elements from the interlayer. This is similar to the previous study of the oxidation behavior of ZrB2-SiC composite. The SiC depleted region can be presented in the low pO2 atmospheric condition due to the evaporation of SiO [34,35]. The evolved SiO gas phase might be evacuated through the melted oxide phases and fully oxidized SiO2 melted phases were dispersed along HfO2 grain boundary and filled the crater of surface layer (region 1 from Fig. 8(a)) to form flow/diffusion-blocking barrier. The existence of partially oxidized hafnium phase is similar to the results from previously reported research for the oxidation behavior of HfC ceramic exposed to a CO2 laser source [36,37]. The oxidation process induced by the oxyacetylene flame is considered to include various thermos-chemical, physical, and mechanical behaviors due to the very high temperature and pressure resulting from the velocity of the combusting gases. Diverse oxidation behaviors at the surface of each sample would be involved during oxyacetylene testing. The reactions possible during the ablation test under the oxyacetylene torch flame are presented below (Eq. (1)-(8)). Although more complex reactions relevant to the presence of oxidants such as H2O and CO2 (which would increase the oxidation rate) could also be involved, only the oxidation reactions involving oxygen gas were given for simplicity  Fig. 6. (a) The surface oxidized layer  thickness as a function of  testing time measured  from the cross-sectional SEM images (b) The calculated linear ablation rate obtained by  dividing the thickness of the surface oxidation layers by the test time, and the inset shows  the thickness of the residual matrix.  granules are HfO2 phase. The size of the HfO2 granules was larger than the average grain size of the as-sintered HfC-SiC sample (~400 nm). The instantaneous oxidation of the HfC grains caused the formation of HfO2 granules, and the HfO2 granules grew while floating around in the melted SiO2 matrix under high temperature. It is also possible that the gaseous phases that occurred during the oxidation test flowed through during this liquid phase. The evacuation of the generated gaseous phases from underneath the surface oxidation layer during the oxidation test, and the influx of oxygen from the surface to the matrix passed through the liquid phase and the bubble-like craters can be seen as evidence of this [15]. Oxygen diffuses through the HfO2 grain boundary, liquid SiO2, or the path (pore, crack) created by the degradation of the sample during the ablation test. The SiC of the matrix reacts with diffused oxygen and is oxidized to SiO2, then melted at high temperature, and viscous flow to the surface by capillary force and volume expansion [11,12]. SiO2 which was moved to the surface is exposed to high temperature and gas velocity, and is vaporized or blew away from the surface of sample. This process repeated during the test.  Table 1 Comparison of ablation or oxidation properties of UHTC materials.  Materials  Distance between the nozzle and sample(mm)  C2H2 flux (l/s)  O2 flux (l/ s)  Max. testing T. (°C)  Max. testing time(s)  Linear reaction thickness(μm)  Linear reaction rate (μm/s)  Ref.  C/C-HfC-SiC ZrC coated C/C ZrB2−15 vol%SiC ZrB2−30vol%SiC ZrB2−20vol%SiC HfC−34vol%SiC  10 10 10 10 -  40  0.31 0.31 0.31 0.31 0.26 0.17  0.42 0.42 0.42 0.42 0.53 0.22  2580 2800 2400 2400 2200 2800  120 240 100 100 2400 1800  312 62 310 30 840 264  2.60 0.26 3.10 0.30 0.35 0.14  [13] [28] [31] [31] [18] This study  8509  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  Fig. 7. (a) The surface SEM images after the oxidation test for 20 min and the EDS analysis at the two spots distinguished (spot 1darkened area, spot 2 brighter area). (b) The SEM  images of the HfO2 melted phases in the deficient SiO2 region.  [18].  SiC (s) + O2 (g) → SiO2 (l) + CO2 (g)  SiC (s) + 3/2 O2 (g) → SiO2 (l) + CO (g)  SiC (s) + O2 (g) → SiO (g) + CO (g)  SiO2 (l) → SiO2 (g)  SiO2 (l) → SiO (g) + 1/2 O2 (g)  HfC (s) + O2 (g) → HfO2 (s) + CO2 (g)  HfC (s) + O2 (g) → HfO2 (s) + CO (g)  HfO2 (s) → HfO2 (l)  (1)  (2)  (3)  (4)  (5)  (6)  (7)  (8)  The sequential changes of the HfC-SiC ceramics during the high temperature oxidation test were summarized in Fig. 9 for better understanding of their probable aspects. It is reasonably expected that some of the oxidation reactions above, such as Eqs. 1-3, 6, and 7 occur  instantaneously (immediately) when the oxyacetylene torch flame reaches the surface of the samples. These oxidation reactions include the generation of gas phases. The gases are released through the liquid phases at the surface. Considering the boiling point of SiO2 based on the observed temperature profile, the evaporation of SiO2 (Eq. (3)) would not have severely occurred during the test. The melted (liquid) SiO2 phase, which formed near the matrix, flowed both outward and into the pores through the channels between HfO2 granules by capillary force. This could be seen as evidence that the amount of SiO2 phase varied with respect to the depth of oxidation from the surface. The oxidation of the HfC-SiC matrix underneath the surface oxidized layer would mainly occur through the instantaneously formed gas evacuation channels in the melted SiO2 phase. This is because there would be no boundary owing to the presence of melted phases during the oxyacetylene testing. The channels for both evacuation and gas inflow through the oxidation layer would be lengthened by the presence of the HfO2 granules in the SiO2 melted phases acting as a flow-blocking barrier. Although the surface oxidation layer is degraded, the combination of melted SiO2 phase and HfO2 granules would adhere well to each other, and they  8510  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  Fig. 8. (a) Magnified cross-sectional SEM images of the HfC-SiC sample tested for 20 min (b) OM image, back-scattered SEM image, and subsequent EDS analysis to confirm the presence  of an interlayer between the surface oxidized layer and matrix, and to identify the elements of  its composition.  enhance the rigidity of the surface oxidized layer without flowing through the surface or spalling, which could cause severe ablation behaviors. The HfC-SiC composite can persist in spite of such severe test conditions owing to the existence of this well supported surface layer.  4. Conclusions  The high temperature oxidation behaviors of HfC-SiC ceramics prepared via reactive spark plasma sintering were investigated under  8511  \\x0c', 'Y.-H. Seong et al.  Ceramics International 44 (2018) 8505-8513  Fig. 9. Schematic illustration of the high-temperature oxidation behaviors of HfC-SiC ceramics under oxyacetylene torch flame.  exposure to the flame of an oxyacetylene torch for up to 30 min. The oxidation rate was determined based on the thicknesses of the oxidized surface layer and the residual matrix. The tested ceramics showed superior resistance to oxidation, compared to previously reported UHTCs materials. The composite structures of the HfO2 micro-granules and melted SiO2 phases have crucial roles in the delay of gas flowing from the surface to the interior matrix. An oxygen-deficient hafnium oxide layer was observed near the matrix. The presence of the hafnium oxide layer supports the high temperature oxidation behavior of HfC-SiC ceramics, and follows diffusion-controlled kinetics. This HfC-SiC ceramic is expected to be beneficial as an impregnation material for various kinds of composite matrices such as C-C or C-SiC fiber matrix to improve the ablation and oxidation resistance and to prevent the matrix from the extreme environmental conditions.  Acknowledgements  This work was supported by the R&D Convergence Program of MSIP (Ministry of Science, ICT, and Future Planning) and under framework of the research and development program of the Korea Institute of Energy Research (B8-2416). The authors would like to thank Dr. S. J. Lee in ADD for his valuable comments and theoretical supports.  References  [1] N.P. Bansal, Handbook of ceramic composites, Kluwer Academic Publishers, Boston, 2005. [2] W. Fahrenholtz, Ultra-high temperature ceramics: materials for extreme environment applications, The America Ceramic Society/Wiley, Hoboken, New Jersey, 2014. [3] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. 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},{
  "_id": 70,
  "PDF": "Evaluation of the high temperature performance of HfB2 UHTC particulate filled Cf-C composites.pdf",
  "Text": "['O R I G I N A L A R T I C L E  Evaluation of the high temperature performance of HfB2 UHTC particulate filled Cf/C composites  Anish Paul1,†  | Virtudes Rubio1  |  Jon Binner1  |  Bala Vaidhyanathan2  | Andrew Heaton3  |  Peter Brown3  1School of Metallurgy and Materials,  University of Birmingham, Birmingham,  UK  2Department of Materials, Loughborough  University, Loughborough, UK  3Dstl, Porton Down, Salisbury, UK  Correspondence  Anish Paul  Email: anish.paul@ansaldoenergia.com  Present address  †Ansaldo Energia Switzerland AG, Baden, Switzerland.  Funding information  UK’s Defence Science and Technology  Laboratory, Grant/Award Number:  DSTLX-1000015267  Abstract  Room and high temperature flexural strength and coefficient of  thermal expansion  (CTE)  of HfB2  ultra-high  temperature  ceramic  (UHTC)  particulate  filled Cf/C  composites  are determined along with UHT oxidation behavior. Both room and  high temperature strength of  the composites were found to be broadly comparable  to those of other  thermal protection system materials currently being investigated.  The CTE of  the  composites was measured both along and perpendicular  to the  fiber direction up to 1700°C and the values were found to depend on fiber orien tation by approximately a  factor of 3. Arc-jet  testing of  the UHTC composites  highlighted  the  excellent  ultra-high  temperature  oxidation  performance  of  these  materials.  K E Y W O R D S  borides, ceramic matrix composites, oxidation resistance  1  |  I N T R O D U C T I O N  Ultra-high  temperature  ceramics  (UHTCs)  are  candidate  materials  for  use  as  leading  edges,  control  surfaces,  engine  inlets  and exits,  and engine hot flow path compo nents  in hypersonic vehicles.  In recent years,  these materi als  have  been  extensively  investigated (TPS)1-3  as  innovative  thermal protection components4-6  systems  and  sharp  leading  edge  for  aerospace vehicles  as well  as  for  other  applications where  oxidation  and/or  erosion resis2000°C are  tance  at  temperatures  up  to  and  exceeding  required. The main materials  that  are being researched as  UHTCs  are  the borides  and carbides of  transition metals,  eg HfB2, ZrB2, HfC, and ZrC. They are refractory in nature and have melting temperatures above 3000°C.7-9 The  suitability of  single phase ceramics  is  significantly limited,  however, due to their poor resistance.10 Even with the  thermal  shock  and  oxidation  addition  of  a  second  or  third  ceramic  phase,  such  as SiC or LaB6, the high temperature resistance, thermal required.11 This  these materials  do  not possess  shock  resistance,  or  fracture  toughness  clearly  highlights  the need to adopt  a fiber  reinforced composite  approach.  Carbon  fiber  (Cf)  and  silicon  carbide  fiber  (SiCf) are two obvious choices, provided they can be pro tected at  the application temperatures.  There are a number of  reports in the literature about  the  preparation of continuous fiber reinforced UHTC composites using SiC4,12,13 and C fibers.14-30 Processing methodologies  adopted for  the preparation of UHTC composites precursor infiltration pyrolysis,18-21 chemical deposition,22-24 reactive melt infiltration,25,26 slurry infiltration and pyrolysis,14-17 or a combination of processes.27-29 A  include  and  vapor  number of groups dedicated their composites,30-35  efforts  to prepare  short  fiber  reinforced  the  advantage  being  the  ability  to  apply  the  processing  techniques  developed  for  monolithic UHTC materials. However,  the improvement  in  mechanical properties, especially toughness,  achieved with  the latter class of materials was not significant. Previous studies conducted by the present authors16 com pared the high temperature oxidation performance of a vari ety of Cf-based UHTC composites viz., Cf-ZrB2, Cf-ZrB220 vol% SiC, Cf-ZrB2-20 vol% SiC-10 vol% LaB6, CfHfB2, and Cf-HfC, using an oxyacetylene flame and reported  that  the  best  performance was  observed  for Cf Received: 14 December 2016  |  Accepted: 5 January 2017  DOI: 10.1111/ijac.12659  344  |  © 2017 The American Ceramic Society  wileyonlinelibrary.com/journal/ijac  Int J Appl Ceram Technol. 2017;14:344-353.  \\x0c', 'HfB2. The main focus of the present study was to determine the room and high temperature flexural strength of these  UHTC composites  together with the coefficient of  thermal  expansion (CTE) along and across the fiber direction. Ultra high temperature oxidation tests were carried out using an  arc-jet facility, which is considered as the best ground based  testing technique for evaluating high temperature oxidation  performance. Arc-jets provide conditions that are similar  to  the aero-thermal environment experienced during flight and  hence the results are used to understand the thermal perfor mance of materials and systems under controlled aero-ther mal  heating  conditions.  The  results  have  been  used  to  validate the numerical models of materials and systems that are used as design tools.5 Nevertheless,  there are a number  of differences between arc-jet and flight environments  that  must be accounted for when interpreting the data. For exam ple, surface catalycity can play a more significant role during  arc-jet  testing than in re-entry, because a higher proportion  of the air molecules are dissociated in the former environment.36 Detailed microstructural characterization was carried  out on the post  test  samples and conclusions drawn about  the  advantages  of  incorporating UHTC particles  on  high  temperature performance.  2  |  E X P E R IM E N T A L  2.1 Preparation of HfB2 UHTC particulate filled composites  |  The  composites  used  in  the  current  study were  prepared  utilizing a slurry composed of HfB2 (325 mesh, HC Starck, Karlsruhe, Germany), acetone, and phenolic resin (Cel lobond J2027L, Hexion Specialty Chemicals, B. V., Rotter dam,  the Netherlands). The  ingredients were mixed in the  required ratio,  a  typical  formulation contained 40 g HfB2, 20 g phenolic resin and 12.5 g acetone, and ball milled for  48 hours to achieve a slurry with the required consistency (~10 mPa seconds at 100 s-1  shear  rate). Cf preforms with  2.5D structure were  obtained  from Surface  Transforms  (Surface  Transforms  plc.,  Cheshire,  UK).  Cf  Preforms  which measure  about  180930915 mm were  used  for  preparing UHTC composites for flexural strength and CTE measurements; 52 mm diameter by ~20 mm thick preforms  were  used  for  arc-jet  samples.  Impregnation  of  the  pre forms was  carried  out using  a vacuum-assisted  technique  where  the  preform was  fully  submerged  in  the UHTC  slurry contained in a vacuum chamber. The  chamber was  then  evacuated with  a  vacuum pump  to  facilitate  the  impregnation of  the slurry. Further details on the composite elsewhere.16  processing  can  be  found  Final  densification  was  achieved using five  cycles of  chemical vapor  infiltra tion (CVI) of carbon at Surface Transforms plc. After CVI  densification, flexural strength (140925910 mm) and CTE  (109595 mm) specimens were machined out  from the lar ger  composites. CTE specimens were machined such that  measurements could be made both along and perpendicular  to  the  fiber  direction. Arc-jet  specimens were machined  down to final dimensions of 30 mm diameter95 mm thick ness,  so  that  they  could  be mounted  in  a  carbon-carbon  (CC)  composite  sample holder. Cf/C composites parative measurements were prepared by CVI densification  for  com of  2.5 D carbon  fiber  preforms  at  Surface  Transforms,  without  any  powder  impregnation.  The  bulk  density  of  HfB2 particulate filled Cf/C composites and Cf/C composites were 2.2\\x060.14 and 1.8\\x060.04 g cm-3, respectively. The final porosity of all the composites was around 10%.  2.2 Flexural strength and CTE measurements  |  Room temperature (RT) and high temperature (HT) 4-point  flexural CERAM1  strength  measurements  were  carried  out  at  (Stoke-on-Trent, Staffordshire, UK) according to  Documented  In-House  Methods  R102:1990  and  R101:2002,  respectively.  The  test  setup was  calibrated  before testing began and the deflection measurements were  initiated after  the  test  temperatures were  reached to mini mize  the  influence of  the  thermal  expansion of  the  exten someter. RT strength measurements were conducted in air  whereas HT strength measurements were carried out under  a flowing argon atmosphere. The  strength of  the  compos ites was  determined  using  large,  140925910 mm,  sam ples;  this was essential  to give a true representation of  the  UHTC composite because of  its graded structure. As per  the specification, prior  to HT testing,  the composites were  coated  with  a  commercial  product  known  as  Tipp-Ex  (a slurry of TiO2 in an organic medium intended for use as a paper correction fluid) all over the surface, except where  it came in contact with the loading and support  rollers,  to  minimize  any  oxidation  due  to  the  presence  of  residual  oxygen. A 5 N preload was applied to ensure proper con tact between the  sample  and the  rollers. The  test  rig used  for HT testing,  along with  a Tipp-Ex  coated  test  bar,  is  shown in Figure 1. The HT test parameters are summarized  in Table 1. RT test parameters were similar except  the fact  that  there was no heating or gas flow.  The argon flow rate employed was  sufficient,  in theory,  to replace the atmosphere within the box furnace approxi mately four  times every minute; \\x001. Presence of the furnace employed is not duration at 1400°C was  this  involved a flow rate  of 15 L min  residual oxygen is expected as  a  sealed  system. The  hold  limited to 5 minutes  to minimize  any oxidation at high temperature due  to the presence of  residual oxygen.  1Now Lucideon Ltd.  PAUL ET AL.  |  345  \\x0c', 'The CTE was measured  at  Imperial College London,  UK,  according  to ASTM E831,  using  a Netzsch  402C  dilatometer  (Netzsch-Geratebau  GmbH,  Selb,  Germany)  with a graphite sample holder and pushrod. Samples were 10°C min 1700°C  heated  at  \\x001  from room temperature  to  under helium atmosphere while recording the displacement  of  the pushrod. By calibrating the expansion of  the set-up  with  a  graphite  sample  of  known  thermal  expansion,  the  displacement of  the pushrod was converted in actual  ther mal  expansion  data for the sample. ~1.5 N on  Since  the  pushrod  exerted  a  force  of  the  sample  to  ensure  that  good contact was maintained, data collected at  the highest  temperatures should be treated with caution as compressive  creep might  have  counter-acted  the  thermal  expansion  of  the sample. Specimens were prepared, so that  the CTE val ues  could  be measured  both  along  and  across  the  fiber  directions.  It  is customary to describe the thermal expansion using  Equation 1.  LðT Þ \\x00 LðTref Þ LðTref Þ  ¼ aavg \\x01 DT  (1)  where aavg from Tref to T; L(T)is  is  the average coefficient of  thermal expansion  the  length of  sample at  temperature  T; L(Tref)  is  the  length of  sample at  reference temperature  Tref; DT=T\\x00Tref.  2.3  |  Arc-jet  testing  Arc-jet  tests of  the samples were carried out at  the German L€uft Aerospace  Centre  (Deutsches  Zentrum  fur  und  Raumfahrt, DLR, Cologne, Germany). One UHTC sample  (AJ5-20) was ~20 seconds  tested  at  a  heat  flux  of  5 MW m  \\x002  for  whereas a second sample (AJ10-10) was \\x002 for ~10 seconds. The test parameters in Table 2. The front face temperature  tested at 10 MW m  are  summarized  during testing was measured using a  two color pyrometer Dr. Georg Maurer GmbH— calibrated from 900 to 3000°C  (Dr. Maurer  QKTR1485,  Optoelektronik, Germany)  and a spectral pyrometer (Dr. Maurer KTR1485-Z, Dr. Georg Maurer GmbH—Optoelektronik, Germany), tive at 1 lm and calibrated between 900 and 3000°C. The surface and cross-sectional microstructures  sensi and  chemical compositions of  the arc-jet  samples were studied  using  field  emission  gun  scanning  electron microscopy  (FEGSEM, Leo  1530VP, LEO, Elektronenskopie GmbH,  Oberkochen, Germany) and energy dispersive spectroscopy  (EDS, EDAX Inc., Mahwah, NJ).  3  |  R E S U L T S A N D D I S C U S S I O N  3.1  |  Flexural strength measurement  The stress-strain curves  for  the Cf/C and HfB2 UHTC particulate filled Cf/C composites after RT and HT testing are given in Figure 2 and Table 3 summarizes the flexural  strength data. The  alumina  rollers  failed on at  least  three  occasions during the HT testing,  resulting in step changes  in the  stress-strain curves. This  can be identified from the  stress-strain plots of UHTC-HT5, CC-HT1,  and CC-HT2.  One of  these samples (CC-HT1) was retested and it yielded  a much lower strength of 85.01 MPa. All other composites  deformed  and  did  not  show any  sign  of  obvious  failure.  The Tipp-Ex applied on the  surface  formed a pale yellow  substance, which was identified to be mainly TiO2. A white layer on the surface of the UHTC composite after HT  F I G U R E 1  4-point bend test rig at CERAM. A test bar coated with Tipp-Exâ can also be seen [Color figure can be viewed at  wileyonlinelibrary.com]  T A B L E 1  HT 4-point bend test parameters used at CERAM  Parameter  Value  Test  temperature, °C  1400  Heating rate, °C min  \\x001  50  Hold duration at 1400°C min  \\x001  5  Argon flow rate, L min  \\x001  15  Initial  load, N  5  Cross-head speed, mm min  \\x001  0.5  Support span, mm  80  Loading span, mm  40  T A B L E 2  Arc-jet  test parameters  Parameter  Value  Sample  AJ5-20  AJ10-10  Test duration, s  20.1  10.6  Heat flux, MW m  \\x002  5.1  10.1  Distance from the nozzle exit, mm  160  100  Peak measured temperature, °C  2400  2650  Specific gas enthalpy, MJ kg  \\x001  15.9  Nozzle configuration  50 mm exit diameter  Test gas or atmosphere  Air  346  |  PAUL ET AL.  \\x0c', 'strength  testing was  characterized  using  EDS  (data  not  shown) and found to be HfO2. The oxidation of Ex and HfB2 confirmed the presence of residual oxygen at the test temperature.  the Tipp The Cf/C composites displayed a higher deformation at room temperature compared to the UHTC composites. No  brittle failure was observed at RT or HT, but  rather a small  amount  of  deformation was  observed.  Cf/C to testing, but  composites  were also coated with Tipp-EX prior  it  fell  off completely during the test, possibly due to the degrada tion of  the surface carbon fibers. There was negligible mass  change for the UHTC composites after HT testing, but Cf/C composites had ~12% mass of the test bars at elevated temperatures.  the  loss  indicating oxidation  The average RT 111\\x0620 MPa  strength  of UHTC composites was  and  that  of  Cf/C  composites  was  132\\x0628 MPa. The average HT strength of UHTC compos103\\x0625 MPa ites was and that of Cf/C composites was 126\\x0610 MPa. The RT strength values reported in the literature for UHTC composites include 107 MPa19 or 150-170 MPa12 for Cf-ZrC; 237 MPa for Cf/ZrB2-SiC;28 for Cf-HfC22 and ~100 or 162 MPa for a functionally graded Cf/HfB2-SiC composite,37 ing on whether the SiC or HfB2 not valid to make direct comparisons as the properties of a  25 MPa  the latter values depend side was  in tension.  It  is  composite depend on the fiber volume  fraction, fiber  sur face  treatment, fiber orientation,  amount of porosity,  type  of carbon deposit, processing temperature, and the type and  amount of fillers.  Considering the error bars,  it  is  reasonable to conclude  that  there was no decrease  in the  average 1400°C,  strength of  the  UHTC  and  Cf/C  composites  at  though  it  is  F I G U R E 2  Stress-strain curves after flexural strength testing.  (A) CC at RT,  (B) CC at HT,  (C) UHTC at RT, and (D) UHTC at HT  T A B L E 3  RT and HT strength of UHTC and CC composites  UHTC composites  CC composites  RT  HT  RT  HT  Sample  Max. Str, MPa  Sample  Max. Str, MPa  Sample  Max. Str, MPa  Sample  Max. Str, MPa  UHTC-RT1  139.69  UHTC-HT1  124.96  CC-RT1  142.70  CC-HT1  130.97a  UHTC-RT2  116.98  UHTC-HT2  124.10  CC-RT2  175.59  CC-HT2  108.64a  UHTC-RT3  111.84  UHTC-HT3  119.32  CC-RT3  86.97  CC-HT3  120.88  UHTC-RT4  75.23  UHTC-HT4  85.51  CC-RT4  126.41  CC-HT4  131.26  UHTC-RT5  113.54  UHTC-HT5  62.67a  CC-RT5  129.15  CC-HT5  139.45  RT, room temperature; HT, high temperature; UHTC, ultra-high temperature ceramic; CC, carbon-carbon. aIndicates the failure of  the alumina rollers during testing.  PAUL ET AL.  |  347  \\x0c', 'difficult  to make any statistical conclusions because of  the  failure of  the rollers  in some instances and the partial oxi dation  of  the  composites.  Cf/C strength compared to UHTC composites  composites  showed  a  slightly higher  at  both room and high temperature. This is not  that surprising  as  the  addition of UHTC powder was  expected to reduce  the overall  strength of  the composites by forcing apart  the  tows slightly as the UHTC powder penetrated.  The difficulties associated with the failure of  the support  rollers at high temperature along with a need for  improved  atmosphere  control need to be  addressed in the  future  to  improve the accuracy of measurements.  3.2  |  CTE measurements  The change in length of  the samples with temperature for  the UHTC and Cf/C composites shown in Figure 3A. All samples showed a large expansion around 1000°C, followed by shrinkage around 1250°C.  from the initial heating is  In  one case,  the shrinkage was so strong that after cooling, a shrinkage of about 50 lm on a 10 mm sample  permanent  was  recorded. As  this variation was  also observed for  the  Cf/C samples, allowing the CVI-deposited carbon and/or HfB2 powder undergo some rearrangement at this temperature. It has  it  is  assumed  that  the  carbon  fibers were  to  been reported that  for carbon materials,  the thermal expan sion  in  any  direction  is  equal  to  the  sum of  crystallite  expansions  resolved  in  that  direction  but  a  proportion  of  that  is  accommodated by internal adjustments. The degree  of accommodation is primarily dependent on the preferred  orientation of  the crystallites with a secondary dependence  on the apparent density of carbon. The presence of  submi croscopic porosity is dence;37  responsible for  this  secondary depen the Cf/C and UHTC composites study had a porosity of around 10%. It is also worth noting  in  the  present  that  the  rapid  change  in  dimensions was  observed  above  the  temperature  employed for CVI of  carbon,  the highest  seen by the  sample during processing, prior  to CTE mea surements. This also suggested that  the processing tempera ture may not have been sufficient  to produce materials that  were stable at high temperatures. As a result of  these varia tions, a second round of measurements was also carried out  for  the  same  samples  and  the  results  are  shown  in Fig ure 3B. This  run  produced  rather  smooth  curves without  much change  in slope  and the  average CTE values  from  these measurements are summarized in Table 4.  The average CTE values of found to be 1.63\\x060.13910 respectively, along and across  the UHTC composites and 4.67\\x060.219 the ply. The cor were  \\x006°C \\x001  10  \\x006°C \\x001,  responding values 2.83\\x060.09910 tively. Type of fiber,  for  the Cf/C composites 4.24\\x060.49910 type of matrix, bonding between the  were  \\x006°C  \\x001  and  \\x006°C  \\x001,  respec fiber and matrix, volume fraction of  the fiber and inter-ply  angle  are  all  factors  that  could influence  the CTE values.  The CTE of Cf along the axis is reported to be negligible \\x001) compared to the value in the radial direc(~0910 \\x001).38 Polymer-derived carbon has a CTE tion (~8910 \\x001 and pyrolytic carbon, which is isotropic, \\x001.39 The CTE has a CTE value in the range 4-6910 \\x001.36 So it can lower CTE of Cf/C and UHTC composites along the ply are mainly due to the lower CTE of  \\x006°C \\x006°C \\x006°C  of 2-4910  \\x006°C \\x006°C  of HfB2 be assumed that  is  reported to be 6.3-7.6910  the  Cf along the axial direction. The contribution of each con stituent phase  to the final CTE can be  estimated provided  the mass  fraction of  each of  the  constituents,  ie Cf, poly mer-derived  carbon  (from the  phenolic  resin),  pyrolytic  carbon (from CVI), HfB2 and submicroscopic porosity are known along with the integrity of the bond between the  fiber and matrix.  F I G U R E 3  Change in length with temperature for CC and UHTC composites.  (A)  Initial  run and (B)  repeated run  T A B L E 4  Coefficient of  thermal expansion values of CC and  UHTC composites  Material  aavg3106°C  \\x001 (25-1700°C)  Cf-HfB2 along the ply  1.63\\x060.13 4.67\\x060.21 2.83\\x060.09 4.24\\x060.49  Cf-HfB2 across the ply  CC along the ply  CC across the ply  CC, carbon-carbon; UHTC, ultra-high temperature ceramic.  348  |  PAUL ET AL.  \\x0c', 'The variation in CTE values along and across  the fiber  direction needs  careful  consideration while designing TPS  components  using UHTC composites. This  variation  can  also be used as a design tool  to fabricate UHTC compos ites with tailored CTE values.  3.3 Arc-jet testing of HfB2 UHTC particulate filled composites  |  Figure 4 shows one of  the Cf-HfB2 the time-temperature profiles during testing are given  samples being tested,  whilst  in Figure 5. Figure 6 compares  the images of  the compos ites before and after the test. AJ5-20 has perature of ~2500°C whereas the sample tested at the higher heat flux, AJ10-10 reached around 2700°C. Melting  seen a peak tem of  the UHTC phase was  not  observed  at  5 MW m  \\x002,  whereas melting was  observed  at 10 MW m  \\x002  indicating  that  the actual  temperature experienced by the sample may  have been higher than the measured value of HfO2 ~2900°C).40 The was uniform whereas the higher velocity jet  (melting point  oxide  layer  formed on AJ5-20  removed some  of  the molten materials from the surface of AJ10-10 during  the test. Both samples survived the rapid heating and main tained  their  integrity  indicating  their  ability  to withstand  ultra-high temperatures and thermal shocks. Combining this  with their  lower density, UHTC composites have an advan tage over UHTC monoliths for UHT applications.  Figure 7  shows  the  surface microstructure  of AJ5-20  after  testing. The surface of  the sample indicated the pres ence of defects, Figure 7A. Necking  of  the  particles was  observed, as shown in Figure 7B. Figure 7C shows an area  near  the  edge  of  the  sample, where  the  surface  layer  became delaminated during the test. This delamination may  have been caused by defects generated during the machin ing of  the  composite  to the  required dimensions,  causing  the fibers underneath the surface layer  to be exposed to the  jet. The  carbon  fibers  underwent  severe  degradation  and  the UHTC particles  showed  partial  oxidation;  they were  not exposed to the jet  for  long enough to undergo complete  oxidation, Figure 7D. Similar partial oxidation behavior was reported for TaC during high temperature testing.41,42  AJ10-10 sample  experienced a higher  temperature  and  heat flux  compared  to AJ5-20,  but  the  test  duration was  shorter. The oxide particles were melted and, on solidifica tion,  formed a dense layer as  shown in Figure 8A. Cracks  were  observed  in  this  layer. The  particles  also  formed  a  protective layer  for  the carbon fibers, Figure 8B. An inter esting observation made on the sample was the degradation  and severe pitting of  the carbon fibers near  the edge of  the  composite, Figure 8C. This  type of damage  is believed to  be due  to the  chemical  attack on the fibers by the highly  reactive gaseous  species  in the  jet,  including monoatomic  oxygen. A cross-section of  the  sample revealed the thickfound to be ~45 lm,  ness of  the surface layer, which was  Figure 8D. The surface cracks observed in Figure 8A were  not propagated to the bulk of  the composite, offering pro tection for  the underlying carbon fibers. study,5  In a previous  arc-jet  testing has been used on  monolithic HfB2-20  vol% SiC ceramics  for much  longer  F I G U R E 4  A picture of one of  the samples being arc-jet  tested,  showing the demanding nature of  the test  [Color figure can be viewed  at wileyonlinelibrary.com]  F I G U R E 5  Time-temperature profile during the arc-jet  testing of  UHTC composites.  (A) AJ5-20 and (B) AJ10-10  PAUL ET AL.  |  349  \\x0c', '350  |  PAUL ET AL.  (A)  (C)  (B)  (D)  F I G U R E 6  Ultra-high temperature ceramic composites before and after arc-jet  testing.  (A) AJ5-20 before test,  (B) AJ5-20 after  test,  (C)  AJ10-10 before test, and (D) AJ10-10 after  test  (A)  (C)  (B)  (D)  F I G U R E 7  Surface microstructure of AJ5-20 after arc-jet  testing.  (A) Surface of  the sample,  (B) higher magnification image showing  necking of oxide particles,  (C)  is an area where the fibers were exposed to the jet, and (D)  is a higher magnification image of  the highlighted  area showing partial oxidation of ultra-high temperature ceramic particles  durations  than in the current work but also at much lower  heat fluxes,  20 minutes  the  formation  of  a  thick,  at 285-350 W cm ~340 lm, removal. The  and  to  due  oxide  layer,  SiO2 formed helped oxygen to diffuse through the material, yielding a SiC depletion layer around ~740 lm thick below  channels  porous  the oxide layer.  \\x002,  and showed  highly  porous  Pressureless  sintered monolithic HfB2 containing 5 vol% MoSi2, have also been arc-jet the stagnation point heat flux was  5-8 MW m  and HfC, both tested,43  for  \\x002  HfB2-MoSi2, 10 MW m  \\x002  which  tested  for  30 seconds,  for  the HfC-MoSi2, which was 4 minutes. The HfB2-based sample was covered 15-20 lm silica-based layer  bubble  and  tested  formation was  the  and  for  by  a  \\x0c', 'observed as a result of  the escape of gaseous by-products.  The HfC-based sample developed a visibly cracked HfO2 layer with no glassy phase visible. A multilayered scale, ~300 lm  thick,  partially  detached  from  the  unreacted  bulk; a similar  layer detachment was observed for Cf-HfC  particulate torch.16  composite when  tested  using  an  oxyacetylene  Severe  oxidation damage has been for Cf/C composites tested by arc jet,44 even at lower temperatures, since active oxidation starts at 800°C. Oxidative consump described  tion of  the carbon fiber/carbon occurs first by transport of  oxygen down cracks,  leading to denuded fibers, and even tual consumption of  the fibers along the cracks. This oxida tion behavior  is not  found for HfB2 ultra-high temperature ceramic particulate Cf/C composite due to the protection offered by the UHTC particles.16  4  |  C O N C L U S I O N S  The room and high temperature flexural  strength and coef ficient  of  thermal  expansion  of HfB2 UHTC particulate filled Cf/C composites have been determined and compared with those of carbon fiber-carbon composites. The Cf/C higher strength than the  composites  showed  a  slightly  UHTC composites at both room and high temperature, but 1400°C was  the reduction <10 MPa,  in  strength  at  relatively  small,  for both groups. There are hardly any reports in  the  literature on the high temperature flexural  strength of  UHTC composites, but  it  can be  concluded that  the high  temperature flexural strength of  the UHTC composites from  the present study is comparable to those of current genera tion TPS materials at  this temperature.  Coefficient of  thermal  expansion measurements  for  the  UHTC composites  revealed  a  large  variation  along  and  across  the ply. The CTE along the fiber direction is  con trolled by the CTE of  the  carbon fiber  in the  axial direc tion; whilst  that perpendicular  is controlled by the CTE of  the  polymer-derived  carbon,  pyrolytic  carbon  and UHTC  particles.  The arc-jet  test  is the first of  its kind reported for slurry  impregnated UHTC composites. Although the test durations  were short,  the samples retained their shape and the surface  erosion was minimal. The UHTC particles formed a protec tive layer at high temperature which was beneficial  for  the  performance of  the composite.  A combination of  low density, good mechanical proper ties, defect and thermal shock resistance and high tempera ture  oxidation  resistance  displayed  by  the HfB2 UHTC from this study clearly  particulate filled Cf/C composites highlighted their potential for hypersonic applications.  It  is  necessary to carry out high temperature  strength measure ments  under  a  completely inert atmosphere (1700°C or higher)  and  at  even  higher  temperatures  to develop a better  understanding of  these materials at  their application temper ature.  It  is  also  essential  to  conduct  arc-jet  testing  for  longer durations.  A C K N OW L E D GM E N T S  The authors thank the UK’s Defence Science and Technology Laboratory (DSTL) for providing the financial support  (A)  (B)  (C)  (D)  F I G U R E 8  Microstructure of AJ10-10 after arc-jet  testing.  (A) Microstructure formed by the melting of ultra-high temperature ceramic  (UHTC) particles,  (B) carbon fiber protected by the UHTC phase,  (C) severe pitting of fibers near  the edge of  the composite, and (D) a cross section revealing the thickness of  the surface layer  PAUL ET AL.  |  351  \\x0c', 'for  this work under contract number DSTLX-1000015267.  Dr. Luc Vandeperre is thanked for his help with measuring  CTE at  Imperial College, London.  R E F E R E N C E S  1. Opeka MM, Talmy IG, Zaykoski selection for 2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci. 2004;39:5887-  JA. Oxidation-based materials  5904.  2. Chamberlain AL,  Fahrenholtz WG, Hilmas GE,  Ellerby DT.  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},{
  "_id": 71,
  "PDF": "Evaluation of ultra-high temperature ceramics for aeropropulsion use.pdf",
  "Text": "['Journal of the European Ceramic Society 22 (2002) 2757-2767  www.elsevier.com/locate/jeurceramsoc  Evaluation of ultra-high temperature ceramics for  aeropropulsion use  Stanley R. Levinea,*, Elizabeth J. Opilab, Michael C. Halbigc, James D. Kisera, Mrityunjay Singhd, Jonathan A. Salema  aNASA Glenn Research Center, 21000 Brookpark Road, MS 106/5#, Cleveland, OH 44135, USA  bCleveland State University, Cleveland, OH 44135, USA  cUS Army Propulsion Directorate, USA  dQSS Group Inc., USA  Received 20 October 2001; received in revised form 20 January 2002; accepted 20 February 2002  Abstract  Among the ultra-high temperature  ceramics  (UHTC) are a group of materials  consisting of zirconium diboride or hafnium  diboride plus silicon carbide, and in some instances, carbon. These materials offer a good combination of properties that make them  candidates for airframe leading edges on sharp-bodied reentry vehicles. These UHTC perform well  in the environment  for such  applications,  i.e. air at  low pressure. The purpose of  this  study was  to examine three of  these materials under conditions more  representative of a propulsion environment,  i.e. higher oxygen partial pressure and total pressure. Results of strength and fracture  toughness measurements, furnace oxidation, and high velocity thermal shock exposures are presented for ZrB2 plus 20 vol.% SiC,  ZrB2 plus 14 vol.% SiC plus 30 vol.% C, and SCS-9a SiC fiber reinforced ZrB2 plus 20 vol.% SiC. The poor oxidation resistance of factor limiting their applicability to propulsion applications. # 2002 Elsevier Science Ltd. All rights  UHTCs is the predominant  reserved.  Keywords: Borides; Composites; Corrosion; Engine components; Thermal shock resistance  1.  Introduction  The high melting points of refractory metal diborides  recognize that stagnation pressures can be greater than  one atmosphere. Some interest has also been shown in  these materials  for  single use propulsion applications.4  coupled with their ability to form refractory oxide scales  The purpose of this study was to examine three of these  give  peratures  temperature  these materials the capability to withstand temin the 1900-2500 \\x0eC range. These ultra-high 1960s.1  ceramics were  the  developed  in  materials under conditions representative of propulsion  environments,  i.e. higher oxygen partial pressure and  total pressure. Relatively long, multiple exposure cycles  Fenter2 provides a comprehensive  review of  the work  are emphasized. Results of strength and fracture tough accomplished in the 1960s and early 1970s. Additions of  ness measurements, furnace oxidation, and high velocity  silicon carbide are used to enhance oxidation resistance  and limit diboride grain growth. Carbon is also some thermal shock exposures are presented for ZrB2 plus 20  vol.% SiC (abbreviated as ZS), ZrB2 plus 14 vol.% SiC  times  used  as  an  additive  to  enhance  thermal  stress  plus 30 vol.% C (ZSC), and SCS-9a SiC fiber reinforced  resistance. These materials offer a good combination of  ZrB2 plus 20 vol.% SiC (ZSS). HfB2 based compositions  properties  that make  them candidates  for  airframe  were not included in the study due to their high cost.  leading edges on sharp-bodied reentry vehicles.3 UHTC  have some potential to perform well  in the environment  for such applications,  i.e. air at low pressure. However,  2. Materials and experimental procedures  for hypersonic flight in the upper atmosphere one must  2.1. Materials  * Corresponding author. Tel.: +1-216-433-3276; fax: +1-216-433 5544.  E-mail address: slevine@grc.nasa.gov (S.R. Levine).  The materials investigated are listed in Table 1. They  were  fabricated by uniaxial hot pressing by Materials  0955-2219/02/$ see front matter # 2002 Elsevier Science Ltd. All rights reserved.  P I I : S 0 9 5 5 2 2 1 9 ( 0 2 ) 0 0 1 4 0 1  \\x0c', '2758  Table 1  Materials  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  Material type  Fabrication  Composition (abbreviated designation)  Monolithic materials  Hot pressed billets \\x18 10 cm diameter by 2.8 cm high  ZrB2 plus 20 v/o SiC, (ZS)  ZrB2 plus 14 v/o SiC plus 30 v/o C, (ZSC)  Composite  Hot pressed plate with uniaxial fiber reinforcement  35 v/o SCS-9a fibers plus ZrB2 plus 20 v/o SiC, (ZSS)  and Machines, Inc., Tucson, Arizona. The powder con stituents were obtained from H.C. Starck (ZrB2  and  SiC), Elkem (silicon), Asbury Graphite Mills  (carbon),  and the SiC fibers  from Textron Specialty Materials  grains appear to be elongated/tabular (approximately 1.5-3 mm thick by 3-11 mm wide/long). ZS was virtually  pore-free, and individual (width 520 mm)  that  large SiC or ZrB2 particles  could act  as  strength-reducing  (SCS-9a fibers).  flaws were not observed in the polished cross  sections.  The ZS and ZSC billet dimensions were nominally 2.8  Significant grain-pullout occurred during the polishing  cm high by 10 cm diameter. ZS had an average density  of the ZSC material as shown in Fig. 2a. This is attrib of  5.57  g/cm3  (100% of  theoretical density),  and the  ZSC had billet densities of 4.47 and 4.54 g/cm3  (98.2  and 99.9% of theoretical density, respectively).  uted to the presence of  the 30 vol.% C, which appar ently  led to weaker bonding within this UHTC. For  comparison, a ZSC fracture surface is shown in Fig. 2b  Specimens were mounted in epoxy and polished, then  to confirm that the material  is close to fully dense.  examined via field emission scanning  electron micro The  composite panel was prepared by  the filament  scopy (FESEM) with a Hitachi S-4700 Field Emission  winding and slurry deposition technique5  followed by  SEM with XEDS  (X-ray  energy  dispersive  spectro hot pressing in a graphite die. The final composite panel  scopy) capability. A representative ZS material micro was  15.1 by  15.1 by  0.53  cm with a density of  3.47  structure  is  shown in Fig.  1. The ZrB2  grains  (gray  phase) are equiaxed, with the majority of the grains ranging from 6 to 12 mm in width. The SiC (dark phase)  g/cm3, while the theoretical density is 4.60 g/cm3 for 35  vol.% fiber  loading. ZSS  composite  specimens were  mounted  in  epoxy  and  polished,  then  examined  via  Fig. 1. Microstructure of ZrB2 plus 20 v/o SiC ultra-high temperature ceramic.  Fig. 2. Microstructure of ZrB2 plus 14 v/o SiC plus 30 v/o C ultra-high temperature ceramic (a) polished section and (b) fracture surface.  \\x0c', 'optical microscopy and FESEM. A representative ZSS  composite microstructure is shown in Fig. 3a. The SCS9a fibers have a diameter of about 80 mm including the  dual  layer  coating.  It appears  that  the  coating on the  surface of  the SCS-9a fibers remained intact during the  hot pressing process. The fiber volume fraction within  the multilayer plies was at  the  target 0.30. However,  thick matrix  layers between plies  reduced the overall  fiber volume fraction to 0.20-0.25. Significant amounts  of porosity can be seen in the matrix (Fig. 3b). Based on  the density of  the billet and a rule of mixtures calcula tion,  the matrix material  is approximately 70% dense.  Longitudinal sections revealed regularly spaced matrix cracks (400-800 mm apart) oriented perpendicular to the  fibers. These  cracks  are  due  to  the matrix  thermal  expansion coefficient being greater than that of the fiber  and thus  leading  to cracking upon cooling  from the  processing temperature, as is observed in other compo sites such as C/SiC.  2.2. Procedures  2.2.1. Mechanical properties  Flexural  strength—test bars  (3 by 4 by 50 mm, with  four long edges beveled) were machined from each billet  of monolithic UHTC material. Test bars (5 by 6 by 50  mm with no beveling of edges) were also machined from  the ZSS billet. Specimens were tested in four-point flex ure (20/40 mm inner/outer spans, silicon carbide fixture,  0.5 mm/min. loading rate) in air at 25, 1127, and 1327 \\x0eC. The load was applied to the samples immedi ately  after  the  test  temperature was  achieved. Four  monolithic  and two composite  specimens were  tested  per  test  condition. Fractography was  performed  on  various tested monolithic specimens to assess the nature  of the critical flaws.  Creep—monolithic  (ZS  and ZSC)  specimens were  tested in four-point flexure (20/40 mm inner/outer spans, silicon carbide fixture) in air at 1127 and 1327 \\x0eC.  Upon reaching the test  temperature, ZS and ZSC sam ples were loaded to half  the average measured strength  of  the material at  the test  temperature. Thus, ZS sam ples were loaded at 250 and 180 MPa at 1127 and 1327 \\x0eC, respectively. ZSC samples were loaded at 100 and 90 MPa at 1127 and 1327 \\x0eC, respectively. Samples  were  held  for  5  h. The  steady  state  creep  rate was  determined. Two to four  runs were attempted per  test  condition.  Fracture  toughness—single  edge  pre-cracked  beam  testing in accordance with ASTM C 14216 was used to  determine  the  fracture  toughness of  the ZS and ZSC  materials. Six tests were conducted per material.  2.2.2. Furnace oxidation  Sample coupons were 2.54 by 1.27 by 0.32 cm. Cou pons were ultrasonically cleaned successively in deter gent, de-ionized water,  acetone  and alcohol prior  to  exposure.  Initial  sample weights  (to  an  accuracy  of  0.00005 g) and dimensions (to an accuracy of 0.001 cm)  were recorded. Three samples were loaded into a slotted  ZrO2 refractory brick. Samples were exposed to tenminute oxidation cycles in stagnant air at 1327 \\x0eC in a  box  furnace with molybdenum disilicide  heating  ele ments (CM, Inc. Rapid Temp Furnace, Bloomfield NJ).  One sample was removed after one cycle, five cycles and  ten cycles. A maximum exposure time of 100 min was  thus achieved. Similar exposures were conducted at 1627 and 1927 \\x0eC using a bottom-loading furnace with  zirconia heating elements (DelTech, Inc., Denver, CO).  Weight change was measured, where possible. Some of  the samples stuck to the sample holder during oxidation  due  to  extensive  glass  formation. X-Ray  diffraction  (XRD) was used to identify oxide phases present after  exposure. After  surface microstructural  analysis  by  SEM and XEDS, samples were cross-sectioned and polished to 1 mm diamond in non-aqueous polishing  media. Water was avoided to preserve any boria that  might be present as an oxidation product. The amount  of surface recession was determined from the difference  between  the  initial  thickness  and  the  thickness  of  unreacted material  that was measured in low magnifi cation  SEM micrographs  of  sample  cross-sections.  Fig. 3. Microstructure of polished sections of ZrB2 plus 20 v/o SiC plus SCS-9a fibers composite showing (a) representative fiber distribution and (b)  matrix porosity.  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  2759  \\x0c', 'Thickness measurements  obtained  from micrographs  were  corrected  based  on  a magnification  reference  standard.  2.2.3. Thermal shock  Thermal shock tests were run by placing 1.27 by 0.32  by  8.31  cm specimens of ZS or ZSC specimens  in a  copper  block  holder  in  front  of  a  hydrogen-oxygen  combustor and nozzle assembly operating at a chamber  pressure of 1.75 MPa. For comparison purposes, a sili con nitride  specimen with similar  test  section dimen sions except  for a width of 1.02 cm in the gage section  was tested. The flame impinged along the length of  the  specimen on the thin edge (thickness) dimension. Initial  5 s  runs were made at an oxygen to hydrogen (O:H)  mixture ratio of 1.7 to achieve a target gas temperature of 1327 \\x0eC. Two runs were made at an O:H mixture  ratio of 2.3 to achieve a target gas temperature of 1627 \\x0eC. The flame was turned from off to full on and  from full on to off. Then the specimens were exposed to  an ambient  temperature nitrogen purge that  impinged  on the cooling specimen. Due to mixing problems in the  combustor,  the target  temperatures were not achieved.  Hot streaks to temperatures much higher than the target  temperature were  present  near  the  periphery  of  the  flame, and the center of  the specimen ran cooler as a  result of oxygen gas flow through a central  igniter tube.  In the earlier tests the specimens were clamped loosely  in the holder at one end. A compliant high temperature  fabric was used to clamp the specimen ends in the later  tests. Visual  observations,  photographs,  some XRD  analysis of  the surface, and metallography documented  test results.  3. Results  3.1. Mechanical properties  Flexural  strength—limited  flexural  testing  in  fast  fracture was performed for the ZS and ZSC materials,  and results are presented in Tables 2 and 3, respectively.  The flexural strength of ZS is comparable to a reported  literature value of 408 MPa at room temperature.3 The  large scatter in our limited number of tests is indicative  of material variability due to poor process control. The  ZSC material has  a  lower  average  strength at  room  temperature, and shows diminished strength at higher  temperature. The 30 vol.% C is  the probable cause of  the reduced strength. The grain pullout observed during  polishing  of  the  samples  indicates weaker  bonding  within the material.  Fractography was performed to assess  the nature of  the critical flaws in various monolithic specimens tested  at room temperature. The fracture origins were usually  difficult  to  locate. Large  clusters of  relatively  coarse  ZrB2 grains surrounding groups or clusters of large SiC  grains were observed within the  fracture mirror  and  fracture was often attributed to these features. The for mer often took on a more or less spherical shape indi cating the possibility of spherical agglomerates that had  not separated from the surrounding matrix during sin tering. The SiC clusters were  in a chain-like  series of  grains within the ZrB2 matrix indicating the possibility  of partially broken down agglomerates. Because of  the  large scatter in the data, no trends with increasing tem perature could be discerned.  Table 2  Mechanical properties of ZS ultra-high temperature ceramic  Four-point flexure tests  Test  temperature  Number  of tests  Average flexural  strength (\\x1b ) (MPa)  Range  (\\x1b ) (MPa)  25 \\x0eC 1127 \\x0eC 1327 \\x0eC  4  391  293-456  4  497  293-666  4  356  227-500  Flexural creep tests (four-point loading)  Test  temperature  Number  of tests  Stress applied  (\\x1bA) (MPa)  Steady state creep rate (s\\x001)  1127 \\x0eC 1327 \\x0eC  2  250  2.6\\x0210\\x0010 2.5\\x0210\\x009  2  180  Fracture toughness (single edge precracked beam (SEPB))  Test  temperature  Number  of tests  Average fracture  toughness (MPa*m1/2)  S.D.  (MPa*m1/2)  25 \\x0eC  6  4.45  \\x06 0.32  Table 3  Mechanical properties of ZSC ultra-high temperature ceramic  Four-point flexure tests  Test  temperature  Number  of tests  Average flexural  strength (\\x1b ) (MPa)  Range  (\\x1b ) (MPa)  25 \\x0eC 1127 \\x0eC 1327 \\x0eC  4  286  270-301  4  206  189-233  4  183  172-205  Flexural creep tests (four-point loading)  Test  temperature  Number  of tests  Stress applied  (\\x1bA) (MPa)  Steady state creep rate (s\\x001)  1127 \\x0eC 1327 \\x0eC  2  100  1.4\\x0210\\x009 1.6\\x0210\\x008  2  90  Fracture toughness [single edge precracked beam (SEPB)]  Test  temperature  Number  of tests  Average fracture  toughness (MPa*m1/2)  S.D.  (MPa*m1/2)  25 \\x0eC  5  4.82  \\x06 0.15  2760  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  \\x0c', 'The  composite  samples  showed  through  thickness  cracks during testing, but did not separate into 2 sepa rate pieces due to crack bridging by the fibers. Delami nations were also observed. Results are summarized in  Table 4.  Creep—minimal deformation was observed for the ZS  and ZSC materials during the 5 h creep tests. Steady state creep rates agreed to within less than a factor of 2  when acceptable tests were run (smooth curve with no  discontinuities). Very low steady state creep rates were  observed 2.6\\x0210\\x0010  for the ZS material (Table 2). Rates of s\\x001 and 2.5\\x0210\\x009 s\\x001 at 250 MPa/1127 \\x0eC \\x0eC,  and  180 MPa/1327  respectively, were  observed.  Total strain (mm/mm) under those follows: 4.8\\x0210\\x005 and 1.5\\x0210\\x004,  conditions was as  respectively.  In spite  of  the presence of 30 v/o C, and the potential  for oxi dation of that C, low ZSC creep rates were observed (Table 3). Rates of 1.4\\x0210\\x009 s\\x001 and 1.6\\x0210\\x008 s\\x001 at 100 MPa/1127 \\x0eC and 90 MPa/1327 \\x0eC,  respectively,  were observed. Total strains (mm/mm) under these conditions were 5.6\\x0210\\x005 and 5.4\\x0210\\x004, respectively.  Fracture toughness—the six valid tests completed on  the ZS material gave KIpb=4.45 \\x06 0.32 MPa ZSC material indicated  a  fracture  toughness  of  p  m. Five valid test results on the  a  fracture  toughness  of  KIpb=4.82 \\x06 0.15 MPa stable crack extension  p  m. All test specimens exhibited  as monitored with  back  face  strain  gages  or  the  actuator  stroke. One  difficulty  encountered in the ZSC testing was measurement of the  crack length:  the  test  specimens  tended to absorb the  dye penetrant use  to mark the  crack fronts. Thus  the  crack length could be measured only if  the pre-crack  plane and the fast  fracture plane were not completely  coplanar,  thereby demarking the precrack front. This  was the case in five of six attempted tests with ZSC.  3.2. Furnace oxidation  Macrographs of  the ZS samples after oxidation are  shown in Fig. 4. Oxide formation is visible on the samtested at 1327 \\x0eC. Extensive glass  ples  formation was  Table 4  Mechanical properties of ZSS ultra-high temperature ceramic  Four-point flexure tests  Test  temperature  Number  of tests  Average flexural  strength (\\x1b ) (MPa)  Elastic modulus  (MPa)  25 \\x0eC 1127 \\x0eC 1327 \\x0eC  2  130  34.5  2  101  33  2  84.5  31  Fig. 4. ZS UHTC oxidized in air for 10-min cycles.  Fig. 5. SEM micrograph of ZS after oxidation in air 1327 \\x0eC.  for 10 min at  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  2761  \\x0c', '2762  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  Fig. 6 shows macrographs of  the ZS sample  cross sections after oxidation. Oxide scales are visible to the unaided eye beginning with exposures at 1627 \\x0eC. The \\x0eC for  exposed  10-min  sample  cycles  1927  10  at  is  almost completely consumed.  SEM and XEDS results for ZS sample cross-sections  are shown in Figs. 7-9. SEM and XEDS results were  After  obtained at 6 kV where sensitivity exposure at 1327 \\x0eC for ten 10-min cycles, the scale was about 30 mm thick and composed of  to boron is high.  oxide  Fig. 6. Cross-sections of ZS oxidized in air for 10-min cycles.  ZrO2 with SiC particles embedded in the scale as shown  in Fig. 7. Fig. 8 shows a ZS sample cross-section after exposure at 1627 \\x0eC for ten 10-minute cycles as well as  the results of an XEDS line scan. The oxide scale was about 150 mm thick. Beginning from the  surface,  the  observed on samples exposed at 1627 \\x0eC. Samples oxidized at 1927 \\x0eC formed an orange oxide after one 10 min  cycle. With  increasing  exposure  time,  the  oxide  scale  became  grayer. At  this  exposure  temperature,  swelling of  the samples occurred with the thickness of  scale was composed of amorphous SiO2  followed by a  layer consisting primarily of ZrO2 in a continuous silica rich glassy phase. No boron was detected in the glassy  phase. A discrete ZrO2/ZrB2 boundary was followed by a SiC depleted region of ZrB2 of about another 100 mm thickness. This SiC depletion layer was observed in  the specimens increasing by up to 80% after ten 10-min  previous work, both the work at high temperatures and  cycles. XRD analyses  show the surface oxidation pro high oxygen partial pressures1 as well as in exposures at  duct is largely monoclinic ZrO2 under all conditions. An  SEM photomicrograph of  dation for one  10-min cycle  the surface of ZS after oxi\\x0eC is  shown in  1327  at  Fig. 5. The large white crystals on the surface are SiO2  lower temperatures, but only in reduced oxygen partial  pressures.7 This SiC depletion was attributed to active  oxidation of the SiC to form SiO(g).7 Fig. 9 shows the ZS sample cross-section after exposure at 1927 \\x0eC for  and the irregularly shaped material  in the background is  ten 10-min cycles. Here,  the oxide  scale was  several  a mixture of SiC, ZrO2 and SiO2.  millimeters thick. The scale was composed of large ZrO2  Fig. 7. SEM micrograph and XEDS analysis of phases in ZS after oxidation in air at 1327 \\x0eC for 10 10-min cycles.  \\x0c', 'S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  2763  Fig. 8. SEM micrograph and XEDS line scan of ZS after oxidation in air at 1627 \\x0eC for 10 10-min cycles.  grains  in a silica-rich glassy phase. No SiC depletion  layer was observed.  Fig.  10  shows  the  surface  appearance of  the  three  materials examined in this study after oxidation exposures of 1, 5, and 10 10-min cycles at 1327 \\x0eC. The most  compact and homogeneous  scale is  formed on the ZS  formulation. This  is  supported by  the weight  change  data for ZS and ZSC shown in Fig. 11. Weight increases  are due to the formation of ZrO2. ZSC oxidizes more  rapidly than ZS due to the presence of porosity resulting  from the oxidation of C and scale disruption due to CO  formation. The weight  loss  for  the SCS-9a fiber  rein forced material is in large measure due to the loss of the \\x18 33 mm carbon cores of  the fibers. Fig. 12 illustrates  Fig. 9. SEM micrograph and the results of XEDS analysis after oxidation in air at 1927 \\x0eC for 10 10-min cycles.  for ZS  Fig. 10. Comparison of UHTC oxidized in air at 1327 \\x0eC for 10-min cycles.  \\x0c', '2764  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  the near parabolic oxidation kinetics and ZSC at 1327 \\x0eC.  exhibited by ZS  tion. Weight change data in Fig. 14 shows near para bolic  behavior  for  ZS,  deviation  from  parabolic  Fig. 13 shows the surface appearance of the ZS, ZSC,  behavior for ZSC due to scale spallation, and a recovery  and ZSS materials after oxidation exposures of 1, 5, and 10 10-min cycles at 1627 \\x0eC. The most  compact and  from the weight  loss of fiber carbon cores in ZSS as a  result of  the rapid kinetics of oxidation of  the porous  homogeneous scale is again formed on the ZS formula matrix to form ZrO2.  Fig.  15  shows  the  surface  appearance of  the  three  materials examined in this study after oxidation exposures of 1, 5, and 10 10-min cycles at 1927 \\x0eC. Extensive  distress in the form of cracks, nodules, and spallation is  evident  in the  scales on all  three materials. ZSS was  severely bloated as a result of oxidation. The weight  change plots  in Fig. 16 are  indicative of  this distress  with parabolic behavior no longer  evident  in ZS, and  large weight losses evident at 5 and 10 cycles in ZSC and  ZSS.  3.3. Thermal shock  Bull  et al.3 used the  following thermal  shock para 1 \\x00 \\x162  \\x01  ¼ 1:8 \\x02 105 W=m  \\x00  \\x0bE  Fig. 11. Oxidation behavior of UHTC at 1327 \\x0eC in air.  meter (TSP) for ZS:  TSP ¼ 24k\\x1bf  Here k is thermal conductivity, \\x1b f \\x16 is Poisson’s ratio, \\x0b is coefficient of  is flexural strength,  thermal expan sion, and E is elastic modulus. The TSP value for  the  AS-800 silicon nitride reference material used in our study is 8.6\\x02105 based on reported property values.8  This  is  indicative  of  better  thermal  shock  resistance  from AS-800 compared to ZS.  Results  of  thermal  shock  tests  are  summarized  in  Fig. 17 and Table 5. One of two ZS samples survived 5 s  exposures at O:H of 1.7, as did a ZSC and the AS-800  sample. Other  tests of AS-800 resulted in survival of  hundreds  of  cycles  (A.E.  Eckel,  private  communi Fig. 12. Parabolic oxidation behavior displayed by ZS and ZSC at 1327 \\x0eC in air.  Fig. 13. Comparison of UHTC oxidized in air at 1627 \\x0eC for 10 min cycles.  \\x0c', 'cation). The AS-800 sample appeared to oxidize  to a  lesser  extent  than  the UHTCs. ZSC oxidized  to  a  greater  extent  than ZS. Longer duration exposures at  O:H of 1.7 resulted in more extensive oxidation and,  in  some  cases,  erosion.  In most  cases  the  samples  frac tured. At the higher O:H of 2.3, samples of each UHTC  material were severely eroded after 180 s of exposure.  As noted after the furnace oxidation exposures, XRD  analysis of  the heat affected zones of  the ZS and ZSC  thermal shock specimens show the oxidation product to  be mainly monoclinic  zirconia.  Some  SiC was  also  detected on specimen F6925 exposed for 210 s at  the  lower O:H ratio. This is in agreement with the furnace oxidation results at 1327 \\x0eC.  In some  cases  the  cubic  zirconia phase was also detected. For  short exposures  where the scales were thin,  the underlying ZrB2 phase  was detected.  Based on the performance of the UHTCs in this test,  thermal shock appears to be a concern in high heat flux  aeroconvective environments such as might be encoun tered in propulsion applications.  4. Discussion  The  three UHTCs  examined in this  study were not  exposed under conditions  for which they appear  to be  best suited, but rather for application in an aeropropul sion environment where oxygen and,  if hydrogen con taining  fuel  is being  combusted, water  vapor partial  pressures are much higher than in reentry to the earth’s  atmosphere from space. The capability of  these materi als  for propulsion applications must be  compared to  mature, available, and commercially used ceramics such  as  silicon nitride, e.g. AS-800,  to put  things  in proper  perspective.  In terms of mechanical properties,  the UHTCs  fall  short in terms of strength and fracture toughness. At about 1300 \\x0eC the creep resistance of ZS as measured  here  (in flexure)  appears  to be  superior  to the  creep  resistance reported for AS-800 (in tension).8 However,  the stress  rupture life for Si3N4 under  stress and tem perature conditions similar to those used in this study is  measured in hundreds of hours.8 ZS could not achieve  such lives due to the oxidation issue. At  their current  state of  early development  it appears  that  reproduci bility of mechanical properties, and thus material relia bility is an issue for UHTCs whereas  the considerable  efforts put  into process development  for  silicon nitride  over the past 30 years have tightened material property  scatter.8  In terms of oxidation resistance, acceptable amounts  of material recession in one hour to thousands of hours,  depending on the specific propulsion application, are on the order of 100-300 mm. This converts to an acceptable  Fig. 14. Oxidation behavior of UHTC at 1627 \\x0eC in air.  Fig. 15. Comparison of the oxidation behavior of UHTC oxidized in air at 1927 \\x0eC for 10-min cycles.  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  2765  \\x0c', '2766  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  range of parabolic recession rate constants, kp of approximately less than or equal to 10\\x001 to 10\\x002 mm2/h  00 ,  for a 1-h application. For a 100-h application an accep00 would be less  than or equal  table range of kp  to the  above values divided by 100. Measured parabolic reces sion rate  constants  in mm2/h for the more oxidation 4.7\\x0210\\x003 \\x0eC, resistant ZS material were at 1327 7.8\\x0210\\x002 at 1627 \\x0eC, and 1.3 at 1927 \\x0eC.9 Thus recession rate constants for ZS are acceptable at 1327 \\x0eC for  Fig. 16. Oxidation behavior of UHTC at 1927 \\x0eC in air.  a 1-h application, but here silicon nitride is a superior material. At 1627 \\x0eC ZS oxidation is marginal for a 1-h  application. But dimensional growth would be an issue.  In a 100-h application ZS cannot be considered at any  temperature.  Prospects  for  the  orders  of magnitude  improvement  in  oxidation  resistance  required  for  UHTC propulsion applications are not good. Zirconia  Fig. 17. Photographs of specimens tested in thermal shock.  Table 5  Thermal shock of UHTCs  Specimen  number  N1382  F6923  F6924  F6925  F6926  F6913  F6914  F6915  Material  O/F  Target temperature (\\x0eC)  Exposure  time (s)  Observations  Si3N4 (AS-800)  ZS  ZS  ZS  ZS  ZSC  ZSC  ZSC  1.7  1.7  1.7  1.7  2.3  1.7  1.7  2.3  1327  1327  1327  1327  1627  1327  1327  1627  5  5  5  210  180  5  75  180  Survived, minimal silica flow along surface  Cracked during cool down  Surviveda, minimal oxidation mostly at hot spot  Crackeda, significant oxidation on surface  Significant material  loss, half of sample broke awaya  Surviveda, minimal oxidation on surface  Surviveda, significant oxidation on surface  Cracked, significant material  loss due to spallationa  a A special high temperature fabric used at the tabs to provide a compliant layer between the sample holder and the specimens.  \\x0c', 'rich scales  (or  for  that matter hafnia rich scales  in the  case of hafnium diboride based materials) provide poor  oxidation protection. This  is due  to the  rapid oxygen  transport and the disruptive phase transformation that  are characteristics of zirconia.  Resistance to water vapor is another issue of concern  for materials  in propulsion applications.  It  is known  that SiO2  formers,  e.g. SiC,  lose  silicon as Si(OH)x(g)  species  at high temperature  in the presence of water  vapor.10 This problem is also a major  issue for  silicon  nitride.11 In addition B2O3 reacts with water vapor12 to  form volatile  species  such as HBO4. Thus one would  expect water vapor  to aggressively attack the UHTCs.  Indeed this is reported to be the case (Q.N. Nguyen and  R.C. Robinson, private communication).  Our cursory examination of thermal shock, both from  a theoretical and experimental viewpoint,  indicated that  the ZS and ZSC UHTCs have less thermal shock resis tance relative to AS-800 silicon nitride.  5. Conclusions  Based on the results of  this  limited study,  the three  examined UHTCs  are not  ready  to be  considered as  aeropropulsion materials  for  any  applications  longer  than  a  few minutes. Current materials  suffer  from  aggressive  oxidation  and moisture  attack. Processing  does not appear  to be under  control as  evidenced by  large scatter in mechanical property data. This property  scatter combined with low fracture toughness  leads  to  poor  resistance  to thermal  shock. The key to raising  their potential for any application is to improve process  and thus material  reproducibility. For  long-term pro pulsion applications major  improvements  in environ mental durability are needed. Given the inherent rapid  oxygen transport  rates and disruptive phase  transfor mation  characteristics  of  zirconia  (and  hafnia),  such  orders  of magnitude  improvements  for  propulsion  applications are unlikely.  Acknowledgements  This work was inspired by the NASA Administrator,  Mr. Daniel Goldin.  References  1. Clougherty, E. V., Pober, R. L. and Kaufman, L., Synthesis of  oxidation resistant metal diboride composites. Trans. Met. Soc.  AIME, 1968, 242, 1077-1082.  2. Fenter,  J. R., Refractory  diborides  as  engineering materials.  SAMPE Quarterly, 1971, 2, 1-15.  3. Bull, J. D., Rasky, D. J. and Karika, C. C., Stability character ization of diboride composites under high velocity atmospheric  flight  conditions.  24th  International  SAMPE  Technical Con ference, 1992, pp. T1092-T1106.  4. Opeka, M. M., Talmy, I. G., Wuchina, E. J., Zaykoski, J. A. and  Causey, S. J., Mechanical,  thermal, and oxidation properties of  refractory hafnium and zirconium compounds. J. Eur. Ceram.  Soc., 1999, 19, 2405-2414.  5. Chawla, K. K. and Chawla, N., Processing of  ceramic-matrix  composites. ASM Handbook, Vol. 21, Composites. ASM Interna tional, Metals Park, OH, 2001, pp. 589-99.  6. ASTM C 1421-99, Standard test method for the determination of  fracture toughness of advanced ceramics at ambient temperatures,  Annual Book of ASTM Standards, V. 15.01, American Society for  Testing and Materials, West Conshohocken, PA, 1999.  7. Tripp, W. C., Davis, H. H. and Graham, H. C., Effect of an SiC  addition on the oxidation of ZrB2. Ceram. Bull., 1973, 52, 612-616.  8.  Pollinger, J. P., Status of  silicon nitride component  fabrication  processes, material properties, and applications. ASME Turbo  Expo 97, Paper 97-GT-321. The American Society of Mechanical  Engineers, New York, 1997.  9. Opila, E. J. and Halbig, M. C., Oxidation of ZrB2-based ultra high  temperature  ceramics. Ceramic Engineering  and  Science  Proceedings, 2001, 22(3), 221-228.  10. Robinson, R. C. and Smialek, J. L., SiC recession caused by SiO2  scale volatility under combustion conditions: I, experimental results  and empirical model. J. Am. Ceram. Soc., 1999, 82, 1817-1825.  11.  Jacobson, N. S., Opila, E. J., Fox, D. S. and Smialek, J. L., Oxi dation and corrosion of  silicon-based ceramics and composites.  High Temperature Corrosion and Protection of Materials 4, Pts 1  And 2, 1997, 251-252, 817-831.  12.  Jacobson, N. S., Farmer, S., Moore, A. and Sayir, H., High temperature  oxidation  of  boron  nitride:  I, monolithic  boron  nitride. J. Am. Ceram. Soc., 1999, 82, 393-491.  S.R. Levine et al. / Journal of the European Ceramic Society 22 (2002) 2757-2767  2767  \\x0c']"
},{
  "_id": 72,
  "PDF": "Evolution of structure during the oxidation of zirconiumdiboride–silicon carbide in air up to 1500.pdf",
  "Text": "['Journal of the European Ceramic Society 27 (2007) 2495-2501  Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500 C     ∗  Alireza Rezaie 1 , William G. Fahrenholtz  , Gregory E. Hilmas  Department of Materials Science and Engineering, 222 McNutt Hall, University of Missouri Rolla, Rolla, MO 65409, United States  Received 22 June 2006; received in revised form 25 September 2006; accepted 7 October 2006  Available online 29 November 2006  Abstract  The structures that developed as dense ZrB2 -SiC ceramics were heated to 1500 C in air were characterized using scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS) and X-ray diffraction. The oxidation behavior was also studied using thermal gravimetric analysis (TGA). Below 1200 C, a protective B2O3 -rich scale was observed on the surface. At 1200 C and above, the B2O3 evaporated and the SiO2 -rich scale that formed was stable up to at least 1500 C. Beneath the surface, layers that were rich in zirconium oxide, and from which the silicon carbide had been partially depleted, were observed. The observations were consistent with the oxidation sequence recorded by thermal gravimetric analysis. © 2006 Elsevier Ltd. All rights reserved.              Keywords: Borides; Composites; Oxidation  1.  Introduction     Ceramic compounds such as ZrB2 , ZrC, TaC, HfB2 , HfC and HfN belong to a group of materials known as ultra high temperature ceramics (UHTCs). Interest in UHTCs has increased substantially in recent years due to growing interest in hypersonic vehicles and re-usable atmospheric re-entry vehicles.1-11 For these vehicles, materials that are resistant to oxidation at 1500 C and above are needed for a variety of components such as nose cones, wing leading edges and engine cowls.4 Currently, UHTCs are among the candidates for these applications as well as other applications that require stability in extreme environments.12-14 As a family of compounds, UHTCs have high melting temperatures (>3000 C) and they maintain their strength at elevated temperatures. With a theoretical density of 6.09 g/cm3 , ZrB2 has the lowest density of the UHTCs,15 which is a desired property for aerospace applications. In addition, ZrB2 has a high thermal conductivity (65-135 W/m K) and has been reported to exhibit excellent thermal shock resistance.15 When ZrB2-SiC is exposed to oxidizing environments at high temperatures, it oxidizes.16,17 Several authors have reported     ∗  Corresponding author. Tel.: +1 573 341 6343; fax: +1 573 341 6934.  E-mail address: billf@umr.edu (W.G. Fahrenholtz). 1 Present address: Vesuvius, Beaver Falls, PA, United States.  0955-2219/$ - see front matter © 2006 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2006.10.012        that oxidation of ZrB2-SiC at 1500 C in air produces a structure that consists of four layers: (1) a continuous silica layer on the surface; (2) a Zr-rich oxidized layer embedded in amorphous silica; (3) a layer of SiC-depleted ZrB2 ; (4) unaffected ZrB2-SiC.16,18,19 A thermodynamic model has been developed to understand the formation of this layered structure.20 In the model, the steady-state response of ZrB2 -SiC to oxidation in air at 1500 C was analyzed with the aid of ZrB2 and SiC volatility diagrams. However, the development of the layered structure on the surface of ZrB2-SiC as it is heated in air to 1500 C and the transient structures that evolve during heating have not been investigated in detail. The purpose of this paper is to describe the structures that develop when ZrB2 containing 30 vol% SiC is heated in air. The compositions of the surface oxides that form during oxidation of ZrB2-SiC are characterized for temperatures ranging from room temperature up to 1500 C.        2. Experimental procedure  2.1. Processing  Commercially available ZrB2 (Grade B, H.C. Starck, Newton, MA) with a reported purity of >99% (metals basis) and an averaged particle size of 2 \\u242em was used. The SiC powder (Grade UF-10, H.C. Starck, Newton, MA) was predominantly  \\x0c', '2496  A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501     ␣-SiC, and it had a reported purity of 98.5% and an average particle size of 0.7 \\u242em. Batches containing 70 vol% ZrB2 and 30 vol% SiC were prepared. Powders were attrition milled using ZrO2 milling media (3.5 mm diameter) to reduce par(Model 01-HD, Union Process, Akron, OH) for 2 h in hexane ticle size and promote intimate mixing. Solvent was removed by rotary evaporation (Model Rotavapor R-124, Buchi, Flawil, (27 kPa), and a rotation speed of 150 rpm. Rotary evaporaGermany) at a temperature of 70 C, a vacuum of 200 mmHg tion was utilized to minimize segregation due to differences in the sedimentation rates of the two powders, which have drastically different densities (6.1 g/cm3 for ZrB2 and 3.2 g/cm3 for SiC). Milled powders were hot-pressed (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 1950 C for 45 min at a pressure of 32 MPa in graphite dies lined with graphite foil that was coated with BN. A detailed description of the temperature ramp used to prepare the specimens has been reported previously.21 A heating rate of 10 C/min to the hot pressing temperature was used. A mild vacuum (20 Pa) was maintained up to 1650 C at which time the atmosphere was switched to flowing argon. An infrared thermometer (Model OS 3708, Omega Engineering, Stanford, CT) was used to monitor the die temperature. A uniaxial load of 32 MPa was applied at 1950 C. After holding for 45 min, the furnace was cooled at 20 C/min to room temperature. Billets with a diameter of 40 mm and thickness of 5 mm were produced. Bars with dimensions of 4 mm × 4 mm × 6 mm and 1.5 mm × 2 mm × 10 mm were sliced from the billets and ground to a 6 \\u242em surface finish for furnace oxidation and thermal gravimetric analysis.                2.2. Oxidation     The experimental portion of this study focused on exposing ZrB2-SiC specimens to air at temperatures of 800, 1000, 1200, 1400 and 1500 C. A MoSi2 resistance heated horizontal tube furnace (Model 0000543, Rapid Temperature Furnace, CM Inc., Bloomfield, NJ) equipped with a high purity alumina tube was used for the oxidation studies. Prior to oxidation, specimens were cleaned in an ultrasonic bath in acetone. Specimens were placed on an alumina plate which was on an alumina D-tube, inserted into the center of the furnace, and leveled. The ends of the furnace were sealed with gas-tight end caps. An atmosphere of flowing air with a volumetric flow rate of 1 cm/s through the tube relative to the specimen (1.8 l/min based on the tube diameter of 6.35 cm) was maintained. Each specimen was heated at C/min to the target temperature and held for 30 min. Specimen temperature was measured with a type B thermocouple that was inserted into the tube and was next to the specimen at an approximate distance of less than 2 cm. After heating to the target temperature, the specimens were cooled to room temperature at 10 C/min. After oxidation, the specimens that had been oxidized at 800 and 1000 C were kept in a sealed container, which was protected from ambient moisture to prevent hydration of B2O3 . The oxidation behavior of ZrB2 -30 vol%SiC was also studied using thermal gravimetric analysis (TGA). The weight change  5           was measured under flowing air at a ramp rate of 5 to 1500 C without an isothermal hold.        C/min up  3. Characterization     The bulk density of each billet was measured using the Archimedes’ technique with deionized water as the immersing medium. The relative density was determined by dividing the bulk density by the theoretical density. The microstructure of each specimen was characterized using scanning electron microscopy (SEM; S-570, Hitachi, Tokyo) along with energy dispersive spectroscopy (EDS; AAT, X-ray Optics, Gainesville, FL) for chemical analysis. For microstructural analysis, cross sections were cut perpendicular to the top surface of the oxidized bars and then polished to a 0.25 \\u242em finish using diamond abrasives. The specimens oxidized in air at 800 and 1000 C were polished with oil based polishing media rather than water to prevent hydration and removal of any B2O3 that was present. For the specimen heated to 1500 C, grazing incidence X-ray diffraction (GXRD; X’Pert MRD, Panalytical, Almelo, Netherlands) was used to determine the crystalline phases present in the SiC-depleted region underneath the surface SiO2 -rich scale after removing the surface layers by polishing parallel to the original surface. The material removal was monitored using optical microscopy so that the desired region was reached. The incidence angle for GXRD was set to 1 , which less than 200 nm into the resulted in a penetration depth of specimen. The GXRD employed Cu K␣ radiation that was passed through a Ni filter. Grain sizes were evaluated using an imaging software program (ImageJ, U. S. National Institutes of Health, Bethesda, MD) by counting a minimum of 250 grains.        4. Results and discussion  4.1. Density  The hot-pressed billets used to produce bars for oxidation studies had measured bulk densities ranging from 5.13 to 5.18 g/cm3 . Using a rule of mixture calculation, and assuming that the true densities were 6.09 g/cm3 for ZrB2 and 3.21 g/cm3 for SiC, the theoretical density of ZrB2 containing 30 vol% SiC was calculated to be 5.23 g/cm3 . Based on this true density, all hot-pressed billets had relative densities of >98%. Consequently, porosity was not considered to have a significant effect on the oxidation behavior.  4.2. Microstructure at room temperature  A microstructure typical of the ZrB2-SiC specimens used in this investigation is presented in Fig. 1. The darker phase is SiC and it appears to be uniformly dispersed in the lighter ZrB2 matrix. SEM analysis did not reveal any obvious porosity in the microstructure, which supports the results of the density measurements. The average grain size of ZrB2 was 2.5 \\u242em, while the SiC particles had an average diameter of 1.7 \\u242em. The microstructure is similar to those reported previously.22,23  \\x0c', 'A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501  2497     (<0.2 mg/cm2 , which corresponds to a layer <0.5 \\u242em thick based on the mass gain and the densities of B2O3 and ZrO2 ). For any B2O3 that did form, hydration and dissolution could have also led to removal of the thin reaction layer during either polishing or storage despite the steps taken to protect it from the ambient moisture. Oxidation of ZrB2-SiC for 30 min at 1000 C led to the formation of enough B2O3 and ZrO2 so that they could be observed in SEM micrographs (Fig. 2). At this temperature, the surface (1) a layer of B2O3 2 \\u242em thick, structure consisted of: (2) layer 6 \\u242em thick that contained unoxidized SiC and a ZrO2 (3) unaffected ZrB2-SiC. A continuous B2O3 layer was found to form above the ZrO2 layer. This structure may form due to volume expansion upon conversion of ZrB2 to ZrO2 and (300% volume B2O3 expansion based on density calculations) and/or the mutual wetting behavior of the two materials. Because the oxidation of SiC is much slower than that of ZrB2 in this temperature regime, the SiC particles did not oxidize appreciably. As ZrB2 oxidized, SiC particles were embedded in the growing ZrO2 layer (Fig. 2). The composition of the layer labeled ZrO2 + SiC was examined using EDS mapping (Fig. 3), which showed that zirconium and oxygen were present along with silicon, suggesting that the reaction layer was composed of ZrO2 and SiC. Determination of the composition of the outermost layer by EDS was not possible due to its low sensitivity to light elements (i.e. boron). However, EDS analysis reported by other investigators has shown that the outermost layer contained O, but was free of Si and Zr, which is consistent with the presence of B2O3 .25-28 Earlier studies have also concluded that B2O3 was an effective barrier to the transport of oxygen, leading to passive oxidation behavior with parabolic mass gain kinetics,24-28 which is consistent with the structure observed in the current study (Fig. 2). Passive oxidation protection is provided by the continuous molten B2O3 layer that effectively seals the surface and prevents direct exposure of the ZrB2-SiC to air in this temperature regime.28  Fig. 1. SEM image of a polished,  thermally etched cross section showing the  microstructure of ZrB2 containing 30 vol% SiC.  4.3.  Initial response during heating (800-1200     C)     Thermodynamically, both ZrB2 and SiC should oxidize when exposed to air. However, the oxidation rates of both species are negligible below about 800 C. Previous studies have reported that the oxidation of ZrB2 by Reaction (1) is much faster than oxidation of SiC (specific reactions are discussed below) between 800 and 1200 C.24,25 Assuming that oxidation of ZrB2 proceeds stoichiometrically, the reaction should produce molten (melting temperature 450 B2O3 C) and solid ZrO2 . Upon cooling to room temperature, the B2O3 forms an amorphous solid while the ZrO2 is crystalline.26 ZrB2 + 5 2 O2 (g) → ZrO2 + B2O3 (l)  (1)           The amounts of B2O3 and ZrO2 on the surface of the specimen oxidized at 800 C for 30 min were not sufficient to be observed in polished cross sections in the SEM. Thermal gravimetric analysis, which will be discussed later in this article, did detect a small amount of oxidation of ZrB2 -SiC at 800 C     Fig. 2. SEM images at (a) low and (b) high magnification showing a layer of B2O3 2 \\u242em thick and a layer of ZrO2 -SiC 6 \\u242em thick formed on the surface of ZrB2 -SiC after exposure to air at 1000 C for 30 min.     \\x0c', '2498  A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501  Fig. 4. SEM image showing the formation of an outer  layer of SiO2 second layer composed of ZrO2 on the surface of ZrB2 -SiC after exposure to air at 1200 C for 30 min.  and a                 resulted in the formation of a continuous surface layer above another oxide layer (Fig. 4). In this case, the underlying layer was composed of porous ZrO2 . A thin SiO2 -rich layer (<1 \\u242em) covered the underlying material and could, potentially, provide a barrier to oxygen diffusion that may result in passive oxidation protection with parabolic mass gain kinetics. This SiO2 -rich layer is expected to contain some B2O3 during transient heating to 1500 C based on either incomplete evaporation of the B2O3 by Reaction (2) or the continued production of B2O3 beneath the outer scale by Reaction (1). Compositional analysis of the outermost layer using secondary ion mass spectrometry (SIMS) has shown that the B content of the oxide layer after heating to C for 30 min is less than 1 wt%.23 1500 The structure formed at 1400 C in the current study (Fig. 5) is consistent with literature reports that indicate that B2O3 evaporates rapidly at temperatures above 1100 C.28 When the evaporating B2O3 is not replaced, as is the case for monolithic ZrB2 , the effectiveness of the diffusion barrier is reduced since the porous ZrO2 layer alone does not protect the underlying ZrB2 from rapid oxidation. For ZrB2-SiC, the addition of SiC extends the oxidation resistance to higher temperatures by promoting the formation of a borosilicate glass layer on exposed surfaces. Previous studies have reported that a SiO2 -rich layer provides passive oxidation protection with parabolic mass gain kinetics, reducing the oxidation rate compared to pure ZrB2 at temperatures above 1200 C.1,24,29 In the present furnace oxidation studies, formation of SiO2 was first observed for the 1200 C specimen (detected by EDS), which is consistent with the previous reports.24 As the temperature approached 1400 layer on the surface increased to 10 \\u242em (Fig. 5). The C, the thickness of the SiO2 thickness of the SiO2 layer, and the underlying ZrO2 -containing layer were not uniform over the specimen surface at this temperature. This may be due to wetting characteristics or other local variations such as composition, surface topology or surface cracks that might enhance the local oxidation rate. The ZrO2 layer was thicker in the areas where the SiO2 layer was thinner, indicating less effective oxidation protection in those regions.           Fig. 3. EDS maps for  (a) Zr,  (b) O and (c) Si  for  the reaction layer  formed  C in air for 30 min showing that a layer that     by oxidizing ZrB2 -SiC at 1000 contains Zr, O and Si is formed.  4.4. Transition structure     As ZrB2-SiC was heated above 1200 C, the composition and structure of the surface layers changed. The dominant chemical processes between 1200 and 1400 C are expected to be the evaporation of B2O3 (Reaction (2)) and oxidation of SiC (Reaction (3)). B2O3 (l) → B2O3 (g) SiC + 3 2 O2 (g) → SiO2 (l) + CO (g)  (3)  (2)     As the temperature approaches 1400 C, the vapor pressure of B2O3 increases substantially,8 leading to its rapid evaporation. In addition, SiC starts to oxidize producing molten SiO2 and gaseous species such as CO in this temperature regime. Like oxidation at lower temperatures, heating to 1200 C for 30 min        \\x0c', 'A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501  2499  Fig. 5. SEM images at (a) low and (b) high magnification of the layered structure formed after exposure of ZrB2 -SiC to air at 1400     C for 30 min.  4.5. Evolution as temperature approaches 1500     C        partially filled with SiO2 that formed through oxidation of SiC at higher temperatures. Although quantitative analysis was not possible by EDS since B is at the limit of the detection capability or by XRD since the B2O3 is amorphous, some B2O3 probably remained dissolved in the SiO2 in this layer, although a previous SIMS investigation showed the amount to be minimal.23 The ZrO2-SiO2 layer was relatively thin (<3 \\u242em). Experiments that have employed thermal cycling appear to promote growth of the ZrO2-SiO2 layer, but no mechanism has been proposed this layer.16 Diffusion of oxyfor the formation or growth of gen molecules or ions through the SiO2 and ZrO2 -SiO2 layers is thought to be the rate determining step in the oxidation of ZrB2-SiC. Materials with coherent SiO2 and/or ZrO2-SiO2 layers exhibit passive oxidation behavior with parabolic mass gain kinetics at 1500 C due to the stability of SiO2 in air in this temperature regime.23,24 A SiC-depleted region, which was located underneath the ZrO2 -SiO2 layer, had a porous structure from which the SiC has been partially or entirely removed (Fig. 6). The morphology of the grains in this region was similar to the original structure before oxidation, except that the SiC had been partially or fully removed by active oxidation. The thickness of the depleted region was 10 \\u242em after heating to 1500 C for 30 min. Grazing incidence X-ray diffraction analysis was used to examine the crystalline phases present in oxidized specimens that had been observed in cross section. After removing the outer SiO2 layer and the ZrO2-SiO2 layer through successive polishing, analysis showed that the SiC-depleted layer was composed of a mixture of ZrO2 and ZrB2 from which some or all of the SiC had been removed (Fig. 7). This result suggests that an oxygen partial pressure (activity) gradient exists across the SiC-depleted layer. In addition, thermodynamic analysis has suggested that the SiO(g) that is generated as SiC is oxidized (Reaction (4) or (5)) is transported across the SiC-depleted layer due to this oxygen partial pressure gradient.20 Since a pO2 gradient is thought to exist across the SiC-depleted layer, an interface separating a layer where ZrO2 is dominant from a layer in which ZrB2 is dominant may be located either: (1) at the interface of the unoxidized ZrB2-SiC and the SiC-depleted layer, (2) in the SiC-depleted layer or (3) at the interface between the SiC-depleted layer and the ZrO2 -SiO2 layer. A more detailed analysis is needed              The structure of the specimen heated in air to 1500 C (Fig. 6) for 30 min was similar to the structure of the specimen exposed to 1400 C, except that the reaction layers were thicker and more uniform after heating to 1500 C. The thickness of the SiO2 -rich layer was 10 \\u242em after heating to 1500 C with a hold of 30 min. At this temperature, the layered structure consisted of: (1) a SiO2 -rich glassy layer; (2) a thin layer of ZrO2-SiO2 ; (3) a layer of ZrO2 and/or ZrB2 from which SiC had been partially depleted; (4) unaffected ZrB2-SiC. This layered structure is similar to the structure reported for ZrB2-SiC exposed to air at 1500 or 1627 C in other studies.16,17 Based on isothermal studies, the SiO2 -rich glassy layer remains protective up to at least 1500 C.1,16,18 Because SiO2 is significantly less volatile than B2O3 at these temperatures (the vapor pressure of B2O3 is 105 times higher than that of SiO2 at 1500 C), the SiO2 -rich layer provides oxidation protection for ZrB2-SiC over a much greater temperature range than the B2O3 does for pure ZrB2 .23 A ZrO2 -SiO2 layer with what appears to be a two phase, interpenetrating microstructure formed beneath the SiO2 layer as the temperature approached 1500 C (Fig. 6). Apparently, the porous ZrO2 that formed initially through oxidation of ZrB2 between 800 and 1200 C was retained, but was covered and                 Fig. 6. SEM image of the layered structure formed after exposure of ZrB2 -SiC to air at 1500 C for 30 min.     \\x0c', '2500  A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501  Fig. 7. Grazing incidence XRD of the SiC-depleted layer formed by oxidation  of ZrB2 -SiC in air at 1500  C.     to identify a distinguishable interface. SiC + O2 (g) → SiO(g) + CO(g) SiO(g) + 1 2 O2 (g) → SiO2 (l)  (4)  (5)     (pO2  Formation of the SiC-depleted layer in ZrB2-SiC specimens exposed to air at 1500 C has been analyzed using volatility diagrams and thermodynamic calculations.20 At  10 intermediate −10 to 10 −15 Pa) oxygen partial pressures that are thought to exist in the SiC-depleted layer, SiC should undergo active oxidation by Reaction (4) or a similar process. The SiO(g) that is formed by active oxidation is transported from the SiC surface (high pSiO(g) and low pO2 ) to the SiO2 layer (low pSiO(g) and high pO2 ) due to the chemical potential gradients of O2 and SiO(g) across the depleted region. At the interface between the SiC-depleted region and the SiO2 -containing layer, the SiO(g) could either oxidize to form additional SiO2 or diffuse into the layer and react with dissolved oxygen to form SiO2 closer to the surface of the outer SiO2 layer. Based on the thermodynamic analysis and the observations reported in this paper, the dominant chemical process in the SiCdepleted layer at 1500 C appears to be the active oxidation of SiC (Reaction (4)) which results in the depletion of SiC.     4.6. Thermal gravimetric analysis (TGA)        To complement the compositional and structural information, the oxidation behavior of the ZrB2-SiC was examined by TGA up to 1500 C with the same heating rate (5 C/min) that was used in the furnace oxidation experiments. The change in the mass as a function of temperature is shown in Fig. 8. The weight started to increase just below 800 C, which corresponds to the temperature at which ZrB2 is reported to begin oxidizing.24,25,28 The weight gain was consistent with SEM analysis that showed a minimal amount of B2O3 formation for specimens heated in air to 800 C (<0.2 mg/cm2 ), but thicker oxide layers at higher temperatures. Between 700 and 1200 at a constant rate (3.3 × 10 −3 mg/cm2   C, the weight increased C). This is consistent           Fig. 8. TGA analysis of ZrB2 -SiC in air up to 1500     C.        with SEM observations that showed formation and growth of ZrO2 and a protective layer of B2O3 (Figs. 2 and 3). As the temperature approached 1200 C the specimen weight decreased of 1.0 × 10 −3 mg/cm2   slightly, with a mass loss rate C between 1215 and 1300 C. The mass loss was attributed to a significant increase in the rate of volatilization of B2O3 .8 Even though the specimen mass decreased in this temperature regime, analysis by EDS verified that SiO2 was present after oxidation at 1200 C whereas no silicon was detected on the surface after oxiC (Fig. 3). Thus, SiO2 is formed at 1200 dation at 1000 C, but the rate of formation must have been less than the rate of evaporation of B2O3 since an overall mass loss was observed by TGA. Above 1300 C the specimen mass increased with a mass gain rate of 1.8 × 10 −3 mg/cm2   C. Based on SEM observations and other analysis, the mass gain was due, primarily, to the formation of SiO2 .              5. Summary           The changes in structure for ZrB2-SiC during heating to 1500 C in air were examined using furnace oxidation followed by SEM/EDS and XRD analysis. Thermal gravimetric analysis was also employed to evaluate the mass change as a function of temperature. Between 800 and 1200 C, oxidation of ZrB2 to ZrO2 and B2O3 was the dominant chemical process. TGA showed that the weight gain started just below 800 C. This resulted in passive oxidation behavior due to the protection provided by the formation of a continuous molten B2O3 layer. Weight gain continued at a constant rate (3.3 × 10 −3 mg/cm2   C) up to 1200 C. At 1200 C, oxidation of SiC was initiated resulting in the formation of SiO2 . In addition, the evaporation of B2O3 became rapid resulting loss recorded by TGA (1.0 × 10 −3 mg/cm2   in a weight The B2O3 was depleted by 1300 C). C, which resulted in mass gain above this temperature due to SiO2 formation. Oxidation of scale increased to a maximum of 10 \\u242em. At 1500 SiC continued up to 1500 C and the thickness of the SiO2 -rich C, a layered structure is formed that consisted of: (1) a continuous SiO2 -rich                 \\x0c', 'A. Rezaie et al. / Journal of the European Ceramic Society 27 (2007) 2495-2501  2501  layer (10 \\u242em); (2) a Zr-rich oxidized layer embedded in amorphous SiO2 (<3 \\u242em); (3) a layer of SiC-depleted ZrB2 and ZrO2 (10 \\u242em); (4) unaffected ZrB2-SiC. The observed increase in the thickness of the outer SiO2 layer required several steps including: (1) the active oxidation of SiC in the SiC-depleted region; (2) transport of SiO(g) across the SiC-depleted region; (3) re-oxidation of SiO(g) to SiO2 . Through the entire temperature range, ZrB2-SiC exhibited passive oxidation behavior in which the diffusion of oxygen through protective molten layers containing B2O3 and/or SiO2 controlled the rate of oxidation.  Acknowledgements  The material presented in this paper is based upon the work supported by the National Science Foundation under grant number DMR-0346800. Additionally, the use of the Advanced Materials Characterization Laboratory at UMR and in particular assistance from Dr. Scott Miller is gratefully acknowledged.  References  1. Levine, S. R., Opila, E. J., Halbig, M. C., Kiser, J. D., Singh, M. and Salem,  11. Valente, T., Marino, G. and Tului, M., Mechanical properties of ceramic  matrix composite for high temperature applications obtained by plasma  spraying.  In Thermal Spray 2004: Advances in Technology and Applica tion, Proceedings of  the International Thermal Spray Conference. ASM  International, Materials Park, OH, USA, 2001, pp. 32-35.  12. Kaji, N., Shikano, H. and Tanaka, I., Development of ZrB2 -graphite protective sleeve for submerged nozzle. Taikabutsu Overseas, 1992, 14, 39-  43.  13. Kinoshita, S., Yoshimasa, Y. and Ono, Y., Application of zirconium boride  materials to waste melting furnace. In Proceedings of International Confer ence on Refractories, UNITECR’03, 2003, pp. 205-208.  14. 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A., Evaluation of ultra high temperature ceramics for aeropropulsion use. J. Eur. Ceram. Soc., 2002, 22, 2757-2767.  19. Opila, E. J. and Halbig, M. C., Oxidation of ZrB2 -SiC. Ceram. Eng. Sci. Proc., 2001, 22, 221-228.  2. Monteverde, F. and Bellosi, A., Development and characterization of metal 20. Fahrenholtz, W. G., Thermodynamic analysis of ZrB2-SiC oxidation: for diboride-based composites toughened with ultra-fine SiC particulates. Solid State Sci., 2005, 7, 622-630.  3. Monteverde, F. and Bellosi, A., The resistance to oxidation of an HfB2 -SiC composite. J. Eur. Ceram. Soc., 2005, 25, 1025-1031.  4. Opeka, M. M., Talmy, I. G. and Zaykoski, J. A., Oxidation-based materials  selection for 2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J. Mater. Sci., 2004, 39, 5887-5904.     5. Van Wie, D. M., Drewry, D. G., King, D. E. and Hudson, C. M., The hyper sonic environment: required operating conditions and design challenges. J. Mater. Sci., 2004, 39, 5915-5924.  6. Chamberlain, A. L., Fahrenholtz, W. G. and Hilmas, G. E., Oxidation of  ZrB2-SiC ceramics under Appl. Trans., 2005, 1, 1-8.  atmospheric  and reentry conditions. Refract.  7. Sciti, D., Brach, M. and Bellosi, A., Oxidation behavior of a pressureless sintered ZrB2 -MoSi2 ceramic composite. J. Mater. Res., 2005, 20, 922-930. 8. Monteverde, F. and Bellosi, A., Oxidation of ZrB2 -based ceramics in dry air. J. Electrochem. Soc., 2003, 150, B552-B559.  9. Nguyen, Q. N., Opila, E. J. and Robinson, R. C., Oxidation of ultra high  temperature  ceramics  in water vapor.  J. Electrochem. Soc., 2004, 151,  B558-B562.  10. Bartuli, C., Valente, T. and Tului, M., High temperature behavior of plasma  28.  sprayed ZrB2 -SiC composite coatings. In Thermal Spray 2001: New Sur faces for a New Millennium, Proceedings of the International Thermal Spray  Conference, Singapore. ASM International, Materials Park, OH, USA, 2001,  pp. 259-262.  mation of a SiC-depleted region. J. Am. Ceram. Soc., in press.  21. Chamberlain, A. L., Fahrenholtz, W. G., Hilmas, G. E. and Ellerby, D. T.,  High strength zirconium diboride-based ceramics. J. Am. Ceram. Soc., 2004, 87, 1170-1172.  22. Rezaie, A. R., Fahrenholtz, W. G. and Hilmas, G. E., Effect of hot pressing  time and temperature on the microstructure and mechanical properties of  ZrB2-SiC. J. Mater. Sci., doi:10.1111/j.1551-2916.2006.01329.x, in press.  23. Rezaie, A. R., Fahrenholtz, W. G. and Hilmas, G. E., Oxidation of zirconium  diboride-silicon carbide at 1500 C a low partial pressure of oxygen. J. Am. Ceram. Soc., 2006, 89(10), 3240-3245.  24. Tripp, W. C., Davis, H. H. and Graham, H. C., Effect of an SiC addition on the oxidation of ZrB2 . Am. Ceram. Soc. Bull., 1973, 52, 612-613. 25. Tripp, W. C. and Graham, H. C., Thermogravimetric study of the oxidation     of ZrB2 in the temperature range of 800-1500 1968, 118, 1195-1199.     C. J. Electrochem. Soc.,  26. Kuriakose, A. K. and Margrave, J. L., The oxidation kinetics of zirconium  diboride and zirconium carbide at high temperatures. J. Electrochem. Soc., 1964, 111, 827-8331.  27. Berkowitz-Mattuck, J. B., High-temperature oxidation. III: Zirconium and hafnium diborides. J. Electrochem. Soc., 1966, 113, 908-914.  Irving, R. J. and Worsley, I. G., Oxidation of titanium diboride and zirconium diboride at high temperatures. J. Less-Common Met., 1968, 16, 102-112.  29. Opeka, M. M., Talmy, I. G., Wuchina, E. J., Zaykoski, J. A. and Causey, S.  J., Mechanical, thermal, and oxidation properties of hafnium and zirconium compounds. J. Eur. Ceram. Soc., 1999, 19, 2405-2414.  \\x0c']"
},{
  "_id": 73,
  "PDF": "Experimental set up for characterization of carbide-based materials inpropulsion environment.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 35 (2015) 1715-1723  Experimental set up for characterization of carbide-based materials in propulsion environment  R. Savino a , G. Festa a , A. Cecere a,∗  , L. Pienti b , D. Sciti b  a University of Naples Federico II, Department of Industrial Engineer, Naples, Italy b National Research Council, Institute of Science and Technology for Ceramics, Faenza, Italy  Received 28 September 2014; received in revised form 16 December 2014; accepted 26 December 2014  Available online 9 January 2015  Abstract  This paper summarizes our achievements in the development of an experimental set up for the characterization of ultra-high-temperature carbidebased material ceramics under conditions representative of a propulsion environment. A segmented converging diverging (de Laval) nozzle was manufactured and  tested  in a  lab-scaled hybrid rocket engine. The converging and diverging sections are manufactured from high  temperature tolerant graphite. A straight section of the graphite nozzle throat was replaced with a low-eroding tantalum carbide composite. The experimental setup enables  to carry  test with different combination of solid propellants and gas oxidizers. The purpose of  this study was  to prepare and  test, under typical operating conditions in the hybrid rocket engine, i.e. very high temperature, high oxygen partial pressure and total pressure, a nozzle throat  insert of a Tantalum Carbide-based composite, fabricated with hot-pressing  technique and characterized  in  terms of strength and fracture toughness, thermal shock, high temperature oxidation behavior. © 2015 Elsevier Ltd. All rights reserved.  Keywords: Ceramic materials; Propulsion environment; Hybrid rocket  1.   Introduction  The  thermal, chemical, and mechanical environments  typical of aero-propulsion applications, such as those characteristic of combustion chambers or of high performance rocket nozzles introduce many problems from  the point of view of materials. These  typical environments are characterized by highly corrosive atmospheres  that may also contain metal additives, with typical flame temperatures even higher than 3000 C. Next generation propellants have become more  energetic  in order  to impart a higher specific impulse to the system, resulting in higher temperatures and pressures  that need  to be contained. These propellants produce very hostile, abrasive environments; existing materials for boost  throat applications have been shown  to erode at unacceptable  rates,  leading  to a  loss  in performance due to throat widening. Implementation of these propellants for boost and thrust applications requires the development of a new     ∗  Corresponding author. Tel.: +39 081 76 823 58.  E-mail address: anselmocecere@hotmail.com (A. Cecere).  http://dx.doi.org/10.1016/j.jeurceramsoc.2014.12.032  0955-2219/© 2015 Elsevier Ltd. All rights reserved.     family of materials providing structural integrity, thermal protection, and lowor near-zero ablation rates above 3000 C. Erosion resistant nozzles that can maintain dimensional stability during firing are  required. The materials used  for  these applications include  refractory metals,  refractory-metal carbides, graphite, plastics.1,2 Certain  ceramics  and fiber-reinforced  classes  of materials demonstrated  superior performances under  specific operating  conditions but  the  choice depends on  the  specific application. For  instance, fully densified refractory-metal nozzles generally are more resistant  to erosion and  thermal-stress cracking  than  the other materials. Graphite performs well with the  least oxidizing propellant but generally  is eroded severely. Some of the refractory-metal carbide nozzles show outstanding erosion  resistance, comparable  to  that of  the best  refractorymetal materials, but generally suffer due to fractures induced by thermal stresses. The interaction of environmental conditions together with the usual requirement that dimensional stability in the nozzle throat can be maintained makes the selection of suitable rocket nozzle materials extremely difficult. This work is focused on the development and testing of monolithic composites based on carbides        \\x0c', '1716   R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723  Fig. 1. Layout of the hybrid rocket engine. HDPE: High Density Polyethlene fuel, Gox: Gaseous Oxygen.        of early transition metals, in particular tantalum carbide.3-5 Ultra High Temperature Ceramics  (UHTC)  such as Tantalum Carbide (TaC) and Hafnium Carbide (HfC) are very good potential candidate materials for use  in propulsive systems. These compounds possess an excellent combination of properties including extremely high melting point  (3950 C and 3928 C,  respectively), high electrical and  thermal conductivity, good  thermal shock resistance and superior ablation resistance compared  to C/C composites. For the above reasons they have been already identified as excellent candidates  for aerospace applications, including also the possibility to develop thermal protection systems  for hypersonic atmospheric  re-entry conditions,  together with other ultra-refractory ceramic composites.6-10 The long-term purpose of this program is to develop thermalshock-resistant  composite materials  systems  that  could  be reliably used with high-flame  temperature  rocket propellants for aerospace propulsion applications.  In concurrent work,11 hot pressing procedures were developed and materials physical properties have been characterized. Preliminary  torch  test have been also carried out at relatively high temperature in significant combustion environments  to  investigate  the materials ablation behavior. Typical conditions encountered  in  real propulsion applications, i.e. the combination of very high temperature, high oxygen partial pressure and  total pressure,  leading  to extreme nozzle thermal and mechanical stresses, can be found in rocket engines. The purposes of  the present  study are  to utilize a  lab-scaled hybrid  rocket engine available at  the University of Naples  to test candidate nozzle materials under the typical operating conditions of a rocket engine. Prior work has shown  that graphite nozzle  throats are eroded and consequently  the physical properties,  in particular  the combustion chamber pressure and  the thrust, change during operation. Therefore, a segmented de Laval nozzle was designed and manufactured with a throat insert of the same ceramic composite material (TaC + 10 vol% MoSi2 ) hot-pressed with the technique described elsewhere.11 The nozzle with  the ceramic  insert was tested comparing  the performance of a similar graphite nozzle throat insert.  2. Experimental  2.1. Device  The Aerospace Propulsion Laboratory of  the University of Naples “Federico  II” was set up primarily  for  the purpose of testing hybrid  rocket  engines.12 Being  equipped with  a  test bench and a general purpose acquisition system, it is extremely versatile  as  it  is  possible  to  easily  adjust  the  experimental apparatus  to  several classes of  tests,  including evaluation of performances of propellants and combustion processes, testing of  sub-components and/or complete power  systems, nozzles, air  intakes,  catalytic  systems, burners,  ignition  and  cooling systems.13-16 The layout of the engine used in the present work, including injector, ignition, single-port cylindrical fuel grain, pre and postcombustion chambers, and expansion nozzle,  is  illustrated  in Fig. 1. The facility was here utilized with  the focus on  test of low erosion  thermal protection materials  in  thrust nozzles.  In order  to evaluate  the performance of hybrid engines,  the small size rocket engine shown  in Fig. 1  is generally equipped with a De Laval nozzle,  to accelerate hot exhaust producing  thrust forces up to 200 N. The pressure in the pre and post-combustion chambers are  in  the  range 5-20 bar. Each  test has a duration that, according to the fuel (Hydroxyl-Terminated Polybutadiene, HTPB, Polyethylene, PE, High Density Polyethylene, HDPE, Paraffin Wax) and to the mass flow rate of the oxidizer (Oxygen or Nitrous Oxide) may be in the order of 20-30 s.17,18 The  test bench  consists of  a  supporting  structure  for  the experimental motor where the thrust is measured by load cells, Tedea-Huntleigh model 1042  (range between 0 and 1000 N). A  pressure  line  reaching  a maximum  value  of  40 bar  supplies  the  oxidizer. Before  the  injection, mass  flow  rate  is evaluated  through gas  temperature and pressure measurements across a Venturi  tube. The pressures  in  the pre-chamber and post-chamber (see Fig. 1) are measured by two capacitive transducers, Setramodel C206 (range between 0 and 35 bar: accuracy: ±0.13% Full Scale). The  acquisition  system  is based on  a software developed  in Labview and National  Instrument PXI  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723   1717  bits. This system,  in addition  to  the management of  the supply lines, allows one to run the tests in a completely automatic way and simultaneously acquire all analog signals. All the pneumatic valves present along the supply lines are controlled and the pressure and mass flow rate are automatically adjusted to predefined set values, coupling the software with a pressure regulator.  2.2. Nozzle design        The  study  in  the present work has been  focused on  the behavior of a converging-diverging (de Laval) nozzle able to be installed in the previously discussed rocked engine. For this purpose, a segmented nozzle made of different materials has been designed and fabricated. The geometry of  the nozzle, with  its main dimension, is illustrated in Fig. 2. The converging element is a  truncated cone with 39 angle, connected  to a  throat with 9.8 mm inner diameter. The final conical element has a diverging angle of 10 and exit diameter of 16.6 mm. The  total  length of the nozzle is 35 mm. During  typical  rocket  tests,  the  nozzle  is  in  graphite  or other ablative materials and  the  throat  is eroded. Preliminary thermal  analyses,  based  on Computational Fluid Dynamics (Fig. 3),  show  that  a  complete nozzle made of TaC  should be extremely  sensitive  to  temperature changes and  therefore extremely  stressed due  to  thermal expansion. A preliminary analysis based on  the Lamé  theory, confirmed by stress computation (Fig. 3c), shows that the maximum equivalent stresses evaluated with the Von Mises method19 are more than one order of magnitude higher than the ultimate and yield material stress (193 MPa) for  the worst case (temperature difference between inner and outer nozzle surface of 1150 K). The reason  is  to be  Fig. 2. Layout De Laval nozzle geometry (dimensions in millimiter).  hardware (PXIe-1082 model) enabling communication between the computer and the equipment. The analog signals from instruments and sensors such as thermocouples, pressure transducers and load cells are sampled with a frequency up to 10 kHz, digitized, processed and stored on a hard disk with a resolution of 16  Fig. 3. Computed temperature distributions (K) in the nozzle for typical operating conditions corresponding to a combustion chamber pressure of 10 bar and oxygen  fuel ratio 3, after 1 s (a) and 20 s (b) from ignition (gas temperature at the entrance of the nozzle of 2280 K); (c) stress computation corresponding to the case (a).  \\x0c', '1718   R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723  Table 1  Operating rocket conditions.  Oxidizer   Fuel   Duration   Chamber pressure   Chamber temperature   Oxygen flow rate   Thrust   O2 HDPE  10 s  6.8 atm  3390 K  27 g/s  57 N  The graphitic parts of the segmented nozzles were machined from a commercial graphite product (Tokai carbon G330). The microstructure of  the ceramic and graphite parts of  the nozzle were analyzed after the test by conventional optical microscopy and scanning electron microscope (FE-SEM, Carl Zeiss Sigma NTS GmbH, Oberkochen, Germany)  and  energy-dispersive spectroscopy (EDS, X-Act, INCA Energy 300, Oxford Instruments, Abingdon, UK)  to understand  the degradation of  the ceramic and graphite parts.  2.4. Experimental set up and operating test conditions  The experimental set up is prepared assembling the propellant grain in the steel container between pre-chamber and the afterchamber (see Fig. 1) and finally the thrust nozzle. In the present study the engine combustion chamber has a length L = 240 mm. The propellant is a cylinder of HDPE with initial inner diameter Di = 14 mm and outer diameter De = 68 mm. Gaseous Oxygen (Gox) is employed. Three  tests have been performed at  the same experimental conditions (see Table 1) with the following nozzles:  (1) monolithic nozzle made of high strength graphite, with a negligible erosion during a single short fire  test. Since  the sonic nozzle throat area remains unchanged during the test stable engine operations and thrust occur during the experiment. With this preliminary test the set-up has been verified and  the values of  the pressure  in  the combustion chamber and of the engine thrust have been evaluated; segmented nozzle with a  throat  insert of molded graphite produced by Tokai Carbon; a segmented nozzle with a throat insert of TaC-based composite material.  (3)   (2)   The assembled test rig with the thrust nozzle and an image of the rocket exhaust plume during a typical test is shown in Fig. 5.  3. Result and discussion  3.1. Microstructure and properties of the starting materials     The TaC-based sample, sintered at 1800 C, is a fully dense material with fine microstructure (mean grain size around 1  \\u242em), and a density of 13.2 g/cm3 . Fig. 6a  is a  typical  fracture surface, which  shows details of  the material bulk. No porosity is observed and secondary phases such as MoSi2 (the sintering aid) and SiC/Si C O based phases are easily  identified.  Fig. 4. Segmented nozzle with converging and diverging elements   in graphite  and throat insert.  ×  found  in  the very high modulus of elasticity of  the TaC up  to 560 GPa and  to  the high  thermal expansion coefficient of  the −6 K −1 ).20 In addition, due to the high density material (6.3   10 of the material (14.5 g/cm3 ) the thermal diffusivity is relatively low compared to conductive materials such as graphite, causing high thermal gradients in transient phases. In  light of  the above  it has been proposed  to provide a segmented nozzle, having  the outer parts,  including converging and diverging conical elements, made of graphite, while  the restricted region around  the  throat  (which  is subject  to higher thermal and mechanical stresses) can be replaced by inserts made of different materials, including bulk ceramics (see Fig. 4). In  this case, being  the part of ceramic material of much smaller size,  thermal stress would be definitely  lower because of the reduced thermal gradients. From preliminary analyses, it is predicted  that  the maximum  temperature differences  in  the nozzle  region are below 150 C and  that  the  thermal  induced stresses be lower than the maximum allowable ones.     2.3. Nozzle fabrication and characterization     A  ceramic  composite with  the  following  composition: TaC + 10 vol% MoSi2 was produced  starting  from  commercial  raw materials  (TaC, Treibacher  Industrie AG, Althofen, Austria), mean particle size 1.1  \\u242em; MoSi2 (<2  \\u242em, Aldrich, Steinbeim, Germany), particle  size  range 0.3-5  \\u242em, oxygen 1 wt%. content  The powder mixture was prepared according to the procedure described in Ref. [11] and hot pressed at 1800 C in low vacuum (100 Pa) using an induction-heated graphite die with an uniaxial pressure of 30-40 MPa; free cooling followed. After sintering, the bulk density was measured by  the Archimedes’ method. The TaC-MoSi2 pellet (with 3 cm diameter and 2.5 cm height) was machined by electrodischarge machining to get the ceramic throat. The bulk microstructure was analyzed  from  thin  sections obtained from the waste parts of the pellet and analyzed by scanning electron microscope (FE-SEM, Carl Zeiss Sigma NTS GmbH, Oberkochen, Germany)  and  energy-dispersive  spectroscopy (EDS, X-Act, INCA Energy 300, Oxford Instruments, Abingdon, UK).  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723   1719  Fig. 5. Assembled test rig (a), (b), (c) with the thrust nozzle and rocket exhaust plume (d) during a typical operative condition.  Fig. 6.   (a) SEM picture of a fracture surface of TaC-based ceramic, showing secondary phases, MoSi2 and SiC. (b) Fracture surface of commercial graphite.  Thermo-mechanical properties  (taken  from Refs.  [11,21]) are summarized  in Table 2. TaC  is a quite strong and stiff material and its thermal conductivity shows the tendency to increase with increasing the temperature at least in the range tested. The  linear CTE  is similar  to other  transition metal carbides.22 The low  fracture  toughness  associated with  the high  stiffness  is probably  the reason for  the poor  thermal shock resistance,  that is however, not notably different  from  those of other UHTC  Table 2  Composition and properties of the TaC-based ceramic and commercial graphite. Properties of the ceramic were measured according to references taken from Refs.  [11,21]. Graphite properties are from the supplier.  Ceramic   Composition (vol%) Bulk density (g/cm3 )/relative density (%)   Young’s modulus (GPa)  Fracture toughness ((MPa m1/2 )   Flexural strength (MPa)   −6 K −1 )  Thermal expansion coefficient (10  Thermal conductivity (W/m K)   Thermal shock resistance (K)  TaC + 10 MoSi2 13.14/99.8   ± ± ±  490   4.7   900    5    0.1    33      C)   6.51 (25-1300     24.1 (20  C), 35 (500     C), 43 (1000     C), 47 (1500     C), 47 (1900     C)   350, by water quenching method   Graphite  Tokai carbon G330  1.8  9.8  -  39.2  4.8  104  -  \\x0c', '1720   R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723  Fig. 7.   (a) Picture of the ceramic throat machined by electro-discharge machining from a pellet with 3 cm diameter and 2.5 cm height, (b) picture of the surface, (c)  SEM of the surface microstructure.  ceramics.23,24 For  the  sake of  comparison, Fig. 6(b)  shows images of the microstructure of the commercial graphite used for machining the segmented nozzle (Tokai carbon G330). Thermomechanical properties of graphite (by the supplier) are also listed in Table 2,  the most striking difference with  the ceramic being the  lower modulus,  lower flexural strength and higher  thermal conductivity.  3.2. Nozzle assembly  Fig. 7a is the TaC ceramic throat and Fig. 7b shows a detail of the surface after machining. The parts of the segmented nozzle and  the assembly  is displayed  in Fig. 8. Basically,  it  is constituted of a graphite main body (1), a ceramic throat (2), an outer graphite ring (3). The ensemble has the same shape and dimensions reported in Fig. 1. For the sake of comparison, a graphite throat was also machined.  CombusƟo n chambe  r  pressu re  Test 1  Test 2  Test 3  )  m  t  a  (  e  r  u  s s  e  r  P  9  8  7  6  5  4  3  2  1  0  0  2  4  6  8  10  12  14  Time (s)  Fig. 9. Comparison between the pressure histories in the combustion chamber  during the 3 tests.  3.3. Rocket engine tests  For all three tests the Oxygen mass flow rate is 27 g/s. Posttest analysis pointed out  that  the  total  fuel consumption after 10 s is about 60 g and that the average oxygen/fuel mass ratio is 4.5.  Fig. 8.   (a) Segmented nozzle, composed of a graphite diverging element   (1),  ceramic   throat made of TaC-based composite material (2), and outer graphite  converging element (3). (b) Ceramic   throat   inserted   into   the graphite body, (c)  Fig. 10. Pictures of   the nozzles after propulsion   tests. Graphite nozzle, and  completed nozzle assembly.  segmented nozzle.            \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723   1721  Fig. 11. Picture of (a) graphite converging element of the segmented nozzle, (b) optical microscopy image of a detail in (a), (c) SEM images of eroded graphite in  (b).  Fig. 12.   (a) Segmented nozzle (after removal of the converging outer element, (b) ceramic throat, (c) details of the ceramic throat showing radial cracks, (d) SEM  image of the interior showing oxide deposits, (e) magnified image of site 1 showing elongated Ta2O5 crystals, (f) magnified image of area 2 in (b), showing Ta2O5 crystals imbedded in silica, (g) EDS analysis of areas in (e) (left) and in (f) (right).  The comparison between  the pressure histories  in  the combustion chamber (post-chamber  in Fig. 1) during  the 3  tests  is illustrated in Fig. 9. Test 3 corresponds to the segmented nozzle with a throat insert of TaC-based material. The time history of the pressure in the combustion chamber during  the  test with  the nozzle made of high strength graphite  (Test 1), shows that the pressure is stable. During the test with the segmented nozzle with a throat insert of molded graphite (Test 2), keeping constant the oxygen flow rate, the pressure reaches a maximum value and then rapidly decreases. The thrust exhibits a similar behavior reaching a maximum value of 58 N and then decreasing during the 10 s test.  \\x0c', '1722   R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723  The average throat erosion rate was measured evaluating the throat diameter before and after  the  test with a digital caliper (accuracy 60 micron) and dividing  the radius variation by  the test duration. During  the first  test with  the monolithic nozzle made of a higher strength graphite, an ablation rate of 0.03 mm/s was measured at  the  throat. The second  test with  the graphite throat insert was characterized by catastrophic erosion due to a combination of chemical reaction and mechanical abrasion; the average erosion rate in this case was 0.28 mm/s, and this explains the sudden pressure decrease as shown in Fig. 9. With  the ceramic  throat  insert  in TaC-based material  (Test 3)  the nozzle remained  in place and no decrease  in  thrust and combustion chamber pressure was noted. After  inspection of the refractory metal carbide immediately after the test, the inner surface of the specimen appeared subject to stress cracking but the inner throat diameter was unchanged.  3.4. Microstructural features after propulsive tests  The visual  appearance of  the nozzles  after propulsion  is shown in Fig. 10, for the monolithic nozzle made of high strength graphite and a segmented nozzle with a  throat  insert of TaCbased composite material. The graphite nozzles both as a unique piece and as segmented assembly showed poor erosion resistance,  in comparison with the refractory metal carbide. Erosion varied from moderate  to catastrophic according to the type of graphite. As an example, the graphite converging element (3)  is displayed  in Fig. 11a-c. As expected, this graphite-based part is heavily damaged compared to the pristine material, due to surface erosion. As  reported  above, with  the TaC-based nozzle  insert no ablation occurred  in  the  throat, Fig. 12. The outer  surface  is unchanged after  the  test,  the material kept  its original color and no visible chemical alteration was observed, although radial cracks were observed (Fig. 12b and c). The ceramic  insert  interior was sectioned and observed by SEM-EDS (Fig. 12d-f). On  the surface of  the  interior,  instead of erosion, an oxidation phenomenon was observed. Following the direction of  the gas flux, deposits consisting of Tantalum oxide (Ta2O5 ) elongated crystals were observed, accompanied by sporadic formation of a glassy silica layer. These phases were the  result of oxidation of TaC and MoSi2 constituent phases. Moreover, residues of C eroded from the frontal part were also noticed. Other regions showed Ta2O5 squared crystals embedded in silica, as confirmed by SEM-EDS images. These surface changes, due  to oxidation, are similar  to  those observed during preliminary  ablation  tests  (Ref.  [11])  even  if much  less pronounced. As for stress cracking, it should be emphasized that there is an important size factor which must be taken into consideration in extrapolating  the  thermal-stress performance of nozzles  in small scale  tests  to full-size applications. The effect of nozzle size on  thermal stresses  is complex and cannot be determined readily. It  is evident  that  thermal stresses  induced  in  the small nozzle insert of this investigation appear to be lower than those that would occur  in a  typical  large nozzle. Accordingly, TaCbased ceramics used in this investigation could not to be suitable  for  large-scale applications. However,  the experimental set up can be utilized for comparative assessment of the relative merits of different materials  in more or  less severe environmental conditions, corresponding,  for  instance,  to  the most oxidizing but  least abrasive propulsion environment, or  to very abrasive exhaust products (e.g. due to aluminized propellant grains).  4. Conclusions  An  experimental  set  up was  prepared  for  characterization of composite ceramic materials under extreme conditions representative of a propulsion environment. A segmented converging/diverging (de Laval) nozzle with an insert of Tantalum Carbide-based material was manufactured and  tested  in a  labscaled hybrid rocket engine and compared with a similar graphite nozzle insert. The  graphite  nozzles  showed  poor  erosion  resistance  in comparison with the refractory metal carbide in the reactive environment.  In particular,  the TaC-based material exhibited high resistance to erosion and maintained dimensional stability during firing, but stress cracking was detected after the test probably explained by high shear stresses during the burn, combined with thermal shocks. Since the experimental setup enables to carry test with different combination of solid propellants and gas oxidizers, including also more  energetic propellants,  further work  is possible  in the  future  to characterize nozzles of different materials when exposed to extreme and very hostile, abrasive environments.  References  1. Hickman R, Mc Kechnie T, Agarwal A. 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Contract No.: HQ0006-05-C-7264.  6. Sciti D, Savino R, Silvestroni L. Aerothermal behaviour of a SiC fibre reinforced ZrB2 sharp component  2012;32(July (8)):1837-45.  in supersonic regime. J Eur Ceram Soc  7. Di Maso A, Savino R, De Stefano Fumo M, Silvestroni L, Sciti D. Arc-Jet  testing on HfB2 -TaSi2 models: effect of  the geometry on  behaviour. Open Aerosp Eng J 2010;3(January (2)):10-9.  the aerothermal  8. Savino R, De Stefano M, Silvestroni L, Sciti D. Arc-jet   testing on HfB2 and HfC-based ultra-high-temperature ceramic materials. J Eur Ceram Soc 2008;28(9):1899-907.  9. Monteverde F, Sciti D, Silvestroni L, Savino R, Esposito A, Carandante  V. Sharp composite UHTC   leading edges for hypersonic applications. In:  Proceedings of the 63rd international astronautical congress. 2012. ISBN:  9781622769797.  \\x0c', 'R. Savino et al. / Journal of the European Ceramic Society 35 (2015) 1715-1723   1723  10. Courtright EL. 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An Italian network to improve  acteristics, processing and applications. Westwood, New   Jersey: Noyes  hybrid  rockets performance:  Astronaut 2014;96:246-60.  the   strategy,   the program,   the   results. Acta  Publications; 1996.  13. Bonifacio S, Festa G, Russo A. Experimental assessment of hydrogen peroxide decomposition in a monopropellant thruster. IJEMCP 2012;10:497-522.  14. Carmicino C, Russo A. Influence of a conical axial injector on hybrid rocket performance. J Propul Power 2006;22(5):984-95.  15. Bonifacio S, Festa G, RussoSorge AM. Novel structured catalysts for hydro gen peroxide decomposition in monopropellant and hybrid rockets. J Propul Power 2013;29(5):1130-7.  16. Carmicino C, Russo Sorge A. Role of injection in hybrid rockets regression rate behaviour. J Propul Power 2005;21(4):606-12.  17. Scaramuzzino   F, Carmicino C,   Festa G, Viviani A, Russo   Sorge A.  21. Sciti D,   Silvestroni L, Guicciardi   S, Dalle   Fabbriche D, Bellosi A.  Processing, mechanical   properties   and   oxidation   behavior   of TaC   HfC  composites  containing  2009;24(6):2056-65.  15 vol%   TaSi2  or MoSi2 .   J Mater   and  Res  22.   Jun CK, Shaffer PTB. Thermal expansion of niobium carbide, hafnium  carbide  and  tantalum  carbide  1971;24(July (3)):323-7.  at high   temperatures.   J Less Comm Met  23. Sciti D, Silvestroni L, Saccone G, Alfano D. Effect of different   sinter ing aids on  thermo-mechanical properties and oxidation of SiC fibers - reinforced ZrB2 composites. Mater Chem Phys 2013;137(January 15 (3)): 834-42.  Fuel   regression-rate characterization on a   lab-scale hybrid   rocket burn 24. Monteverde F, Scatteia L. Resistance to thermal shock and to oxidation of  ing N2O   and   paraffin-based   propellants.   In: Proceedings   of   the   49th  AIAA/ASME/SAE/ASEE joint propulsion conference. 2013.  metal diborides-SiC ceramics for aerospace application. J Am Ceram Soc 2007;90(April (4)):1130-8.  \\x0c']"
},{
  "_id": 74,
  "PDF": "Experimental Verification of Computationally Predicted Preferential Oxidation in Refractory High Entropy Ultra-high Temperature Ceramics.pdf",
  "Text": "['Journal Pre-proof  Part II: Experimental Verification of Computationally Predicted Preferential Oxidation in Refractory High Entropy Ultra-high Temperature Ceramics  Lavina Backman ,  Joshua Gild ,  Jian Luo , Elizabeth J. Opila  PII: DOI: Reference:  S1359-6454(20)30502-4 https://doi.org/10.1016/j.actamat.2020.07.004 AM 16144  To appear in:  Acta Materialia  Received date: Revised date: Accepted date:  24 December 2019 30 June 2020 2 July 2020  Please cite this article as: Lavina Backman , Joshua Gild , Jian Luo , Elizabeth J. Opila , Part II: Experimental Verification of Computationally Predicted Preferential Oxidation in Refractory High Entropy Ultra-high Temperature Ceramics, (2020), doi: https://doi.org/10.1016/j.actamat.2020.07.004  Materialia  Acta  This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.  © 2020 Published by Elsevier Ltd on behalf of Acta Materialia Inc.   \\x0c', 'ijgild@ucsd.edu; iijluo@ucsd.edu; iiiopila@virginia.edu Part II: Experimental Verification of Computationally Predicted Preferential Oxidation in Refractory High Entropy Ultra-high Temperature Ceramics Lavina Backmana,†, Joshua Gildb,i, Jian Luob,ii, Elizabeth J. Opilaa,iii aDepartment of Materials Science and Engineering, University of Virginia, Charlottesville VA 22904 bMaterials Science and Engineering Program, University of California-San Diego, La Jolla, CA †Corresponding author. Postal address: 395 McCormick Road, P.O. Box 400745, Charlottesville VA 22904.  E-mail address: lb2ty@virginia.edu. Phone number: +1 (434) 249 6685 Abstract  Refractory high entropy materials have garnered significant research interest due to their potential ability to fill a need in high temperature structural applications. However, challenges remain with respect to designing for oxidation resistance. A knowledge gap exists with respect to a rigorous understanding of the mechanisms driving oxidation processes unique to high entropy materials. This work provides an experimental complement to a companion publication, which outlines analytical and computational thermodynamic approaches that are envisioned to aid the design of refractory high entropy materials containing group IV (Hf, Zr, Ti) and group V (Ta, Nb) constituents. In this work, (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2) carbide and diboride specimens were exposed at 1700°C in 1% O2 for 5 minutes. Experimental results show good agreement with the computational predictions for the same temperature, despite differences in the overall morphology of the oxidized regions. The carbide formed porous oxides, while the diboride formed a denser external scale. Oxidation products are dominated by group IV oxides, depleting the underlying materials, which were found to consist of primarily group V carbides and borides respectively. The results provide a first look at the oxidation of high entropy                   \\x0c', 'UHTCs at ultra-high temperatures and validate the preferential nature of high entropy material oxidation predicted by the computational approach developed for the study of these new class of materials.  Key words: high temperature oxidation, high entropy alloys, ultra-high temperature ceramics, high entropy carbides, high entropy borides, refractory  1.0 Introduction Ultra-high temperature ceramics (UHTCs), most notably refractory metal carbides, nitrides and borides, may hold the key to advancement in hypersonic flight technology.  UHTCs exhibit melting temperatures exceeding 3000°C, making them appropriate candidates to withstand the extreme temperatures experienced by the leading edges of the vehicle during hypersonic flight.  However, their propensity to react rapidly with oxygen limits their sustained application [1,2].  Researchers have attempted to address this limitation through the modification of well-understood systems. For instance, the addition of SiC to boride UHTCs, such as ZrB2, improves oxidation resistance [2-4] due to the formation of a liquid borosilicate glass layer. However, this layer will shear off at the speeds and temperatures encountered in hypersonic flight (maximum melting temperature for this layer occurs when the scale is pure SiO2 at 1723°C).  Further, the active oxidation of silicon carbide at temperatures higher than 1600°C also precludes the use of silica formers for passivity [5]. High entropy ultra-high temperature ceramics (HE-UHTCs), an emerging class of materials [6-8], are of interest due to the potential to significantly expand the compositional breadth of UHTCs and the range of achievable properties. The design of an HE-UHTC will require not only the ability to form a single-phase solid solution but also to optimize the desired properties including oxidation resistance at ultra-high temperatures (>1700°C), the focus of this paper.                   \\x0c', 'While several studies have been published on the observed oxidation behavior [9,10] of high entropy multi-principal component materials (of which high entropy alloys and ceramics are subsets), a rigorous treatment of the underlying mechanisms that control the oxidation behavior unique to high entropy materials have yet to be established. Prior work has shown that thermodynamic favorability plays a role in the assemblage of phases formed after exposure to the oxidizing environment [9,11], suggesting highly composition dependent oxidation behavior that is thermodynamically driven. A companion publication [12] proposes a method to predict the oxidation products in these complex, solid solution materials, with the goal of using these findings to inform new, oxidation-resistant designs for all non-oxide high entropy materials. The computational work simplifies real materials to the ideal solution model due to a lack of thermodynamic data to consider real systems. An experimental study is therefore needed to validate the predicted assemblage of oxide product and substrate phases. The objective of this work is to apply the previously developed computational approach to experimental oxidation of UHTC systems of interest in order to verify its predictions.   2.0 Methodology Materials chosen for the study are carbides and borides containing group IV and group V elements: Hf, Zr, Ti, Ta and Nb. Hf and Zr carbides and diborides are UHTC candidates that form oxides with the highest melting temperatures [2]. Tantalum has been an extensively studied elemental addition for high temperature UHTCs [13-15], and in the context of this work, has the ability to form solid solutions with Hf and Zr based compounds. Next to group IV oxides, tantalum oxides also have the highest melting temperature among the species under consideration. The (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2) carbide and diboride were also among the first compositions for which single phase, UHTC solid solutions were made [6,16]. The approach outlined in the companion publication [12] was used to                   \\x0c', \"calculate the oxide and resulting substrate composition at 1700°C for the carbide and boride. Note that the partial pressure of oxygen is not defined as an input but is a result from the free energy minimization calculations given the initial condition of a carbide (or boride) solid solution at equilibrium with an oxide solid solution.   The (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C specimen tested in this work was prepared by high energy ball milling and spark plasma sintering (SPS) commercially available powders (Alfa Aesar Haverhill, MA) [16,17]. The (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 specimen was prepared by borocarbothermal reduction [18], followed by SPS consolidation. The sintered samples were machined at Bomas Machine Specialties (Somerville, MA) into 0.5mm thick dogbone specimens shown in Figure 1. Samples cut from the same sintered pucks were mounted, polished and characterized before oxidation via scanning electron microscopy (SEM, FEI Quanta 650, FEI-Thermo Fisher Scientific, Hillsboro, Oregon), energy dispersive spectroscopy (EDS, Oxford Instruments Aztec X-Max^N 150, Concord, MA) and X-ray diffraction (XRD, PANalytical X'Pert Pro MPD or PANalytical Empyrean, Almelo, The Netherlands).  Figure 1: Dogbone specimen configuration. The approximate hot zone, or pyrometer sighting zone (diameter 1.5 mm), is shown in orange. Dimensions shown are in millimeters (mm). Plan View Side View                   \\x0c\", ' Figure 2: Schematic of the resistive heating experimental set-up. Specimen shown is of different configuration for clarity.  These dogbone specimens were then loaded into a resistive heating system, modified from Karlsdottir and Halloran [19], and Shugart and Opila [20], schematic shown in Figure 2. The samples were heated to the desired temperature (nominally 1700°C in this study) through Joule heating. The temperature was chosen as it is the minimum temperature in the ultra-high temperature regime, wherein strategies to promote formation of SiO2 are not expected to be effective. The temperature was controlled by a proportional-integral-derivative (PID) controller and an emissivity correcting one-color pyrometer (Pyrofiber Lab PFL-0865-0790-2500C311, Pyrometer Instrument Company, Ewing Township, NJ). The pyrometer sights on an approximately 1.5 mm diameter zone in the middle of the dogbone specimen (Figure 1). The mean temperatures measured by the pyrometer in this zone during the oxidation test for the carbide was 1715 ± 34.4°C, and the diboride 1700 ± 27°C. The samples were ramped to temperature in flowing (~1L/min) ultra-high purity argon (2 vol-ppm O2 impurity max, Praxair, Danbury, CT), and                   \\x0c', 'certified 1 vol % O2 (balance argon) gas (Praxair, Danbury, CT) was turned on (~1L/min) once the test temperature was achieved. Isothermal oxidation exposures were conducted for 5 minutes from the time the oxidizing gas began to flow. The low partial pressures of oxygen used result in lower oxidation rates allowing for longer exposures and time dependent studies which will be described in a future publication [Backman L., Gild J., Harrington T., Vecchio K., Luo J. and Opila, E.J., manuscript in preparation]. After oxidation, the oxide morphology in both plan view and cross section were characterized using the Everhart Thornley detector (ETD), circular backscatter detector (CBS) and ion conversion and electron (ICE) detector in the SEM, and the composition characterized using EDS. The oxide phases were identified using XRD. The samples were then manually fractured across the hot zone, and the cross-sections were examined in an SEM. Focused ion beam (FIB) milling (Helios UC G4, Thermo Fisher Scientific, Hillsboro, Oregon) was used to obtain samples for imaging and EDS analysis. The schematic shown in Figure 3 shows the location of the FIB lift-out.   Figure 3: Schematic showing the (a) unoxidized dogbone specimen, (b) oxidized dogbone fractured in the middle of the hot zone, with the hot zone shown in yellow and linked to the FIB lift-out with the approximate location in the hot zone. Total lift-out dimensions are approximately 20µm x 30 µm x 1µm before the thinning process.                     \\x0c', '3.0 Results 3.1 Baseline Material The (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C exhibited ~99% relative density, no open porosity and ~4 vol% retained oxide contamination from the starting powders (Figure 4) and was synthesized at University of California, San Diego (synthesis and results from similar samples described in Harrington et al. [16,17]). The single phase (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 was synthesized by consolidating borocarbothermally reduced powders by Gild et al. as described in their publication [18] and exhibited ~99% relative density with ~4 vol% B4C retained from processing and a minor amount (<1 vol%) of rock-salt carbide. The retained oxides in the carbide and the secondary phases in the diboride are independently interspersed. A baseline set of EDS maps of polished cross-sections is provided here in Figure 4 and Figure 5, before oxidation, for reference and comparison to results after oxidation. The distribution of elements is uniform, except for the retained hafnium and zirconium oxides in the otherwise single-phase carbide material. These oxide impurities originate from the surface of the raw material powders used and are difficult to reduce due to their thermodynamic favorability.    Figure 4: Polished cross-sections of (Hf0.2ZrTi0.2Ta0.2Nb0.2)C before exposure to oxygen at ultra-high temperatures. Secondary electron image and EDS maps showing elemental distribution of titanium, tantalum, hafnium, zirconium, niobium and oxygen.                   \\x0c', ' Figure 5: Polished cross-sections of (Hf0.2ZrTi0.2Ta0.2Nb0.2)B2 before exposure to oxygen at ultra-high temperatures. Secondary electron image and EDS maps showing elemental distribution of titanium, tantalum, hafnium, zirconium and niobium.  3.2 Carbides: After Oxidation Figure 6 shows optical and back-scattered electron images of the exposed (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C specimen. The hot zone can be discerned by a dark gray region in the middle of the specimen surrounded by a lighter gray oxide (Figure 6 (c) and (d)). These images highlight the temperature gradients inherent in the experimental set-up, which will be addressed in more detail in future work. This paper focuses on results obtained from the center of the hot zone, well within the boundary of the pyrometer sighting zone.                     \\x0c', ' Figure 6: Plan view optical images showing (Hf0.2ZrTi0.2Ta0.2Nb0.2)C (a) before oxidation, (b-c) after oxidation at a nominal temperature of 1700°C in 1%O2/bal Ar for 5 minutes, and manual fracture along the hot zone, and (d) back-scattered electron image in plan view of the hot zone before fracture. The orange circle is added to (d) to show the estimated size and location of the pyrometer sighting zone.  At least two phases can be discerned on the surface from the microstructure in the hot zone as shown in Figure 7: one that is rich in the group IV elements, and the other rich in Ti, but now with increased Ta and, to a lesser extent, Nb. Figure 8 shows the corresponding indexed X-ray diffraction pattern of oxide formed on (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C. The composition of Phase 1, bright phase in Figure 7, is consistent with a titanate, (Hf,Zr)TiO4 The rounded morphology of the second phase could be indicative of the occurrence of melting. The composition and morphology of Phase 2, dark grey in Figure 7, is a good match for TiTa2O7, which melts at 1674°C [21]. XRD data shown in Figure 8 indicate other possible oxide phases, listed in Table 1. It should be noted that XRD analysis of the oxidized specimens in this study provide an indication of the likely phases present, but multiple overlaps make conclusive analysis difficult. Substition of elements from the same periodic group are also likely, indicated in the third column shown in Table 1. Finally, the carbide was found to oxidize intergranularly                   \\x0c', '(Figure 9 and Figure 10), with the subsurface oxides having different compositions compared to that on the surface; the other phases detected via XRD may be present in these regions.  Figure 7: Plan view images of (Hf0.2ZrTi0.2Ta0.2Nb0.2)C after oxidation at a nominal temperature of 1700°C in 1%O2/bal Ar for 5 minutes. (a) A low-magnification image of the hot zone on one half of a dogbone (orange line indicates approximate boundary of the pyrometer sighting zone), (b) higher magnification, back-scattered electron image obtained at 5kV of the microstructure in center of the pyrometer sighting zone (c) semi-quantitative composition of the two phases (bright and dark gray) apparent in this region from EDS obtained at 10kV.                      \\x0c', '3040506070Other carbide(HfZrTiTaNb)CTiTa2O7m-HfO2HfTiO4+++++++Log Intensity2\\uf071 (degrees)+Zr6Nb2O17 Figure 8: X-ray diffraction spectra taken after oxidation for (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C exposed for 5 mins at a nominal temperature of 1700°C in 1 mol % O2/bal Ar  Table 1: Possible phase matches for the x-ray diffraction pattern (Figure 8) collected from (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C oxidized at a nominal temperature of 1700°C in 1% O2/Ar for 5 mins. PDF numbers or references are indicated in parentheses or brackets, respectively. Indicated Phase Crystal System Possible alternate phases Zr6Nb2O17 (04-011-2890) Orthorhombic (Hf,Zr)6(Ta,Nb)2O17 [22] HfTiO4 (01-080-2977) Orthorhombic (Hf,Zr)TiO4 [23] HfO2 (00-034-0104) Monoclinic (Hf,Zr)O2 [24,25] TiTa2O7 (01-084-8891) Monoclinic Ti(Ta,Nb)2O7 [26] Other carbide Cubic NbC0.92O0.01 (04-022-7393) Nb0.80Ta0.20C (01-085-4222) Zr0.20Ta0.80C (04-002-6875) NbxCx-1 [27] NbC (00-038-1364) NbC0.87 (04-008-3212)                     \\x0c', 'Figure 9 shows the ion conversion and electron (ICE) detector secondary electron (SE) image and EDS maps of oxide scale fracture cross-section, wherein the gas-oxide surface is underneath the carbon layer deposited during the FIB milling process. A white circle is added to an area of interest, which is a partially oxidized carbide grain. The oxygen lean areas correspond with an enrichment of Ta and Nb, whereas the oxygen rich areas, which are porous, are mostly rich in the group IV elements. Figure 9 also shows the Ti-Ta rich oxide phases at the surface (Figure 7), i.e., the gas-oxide interface, where the material is exposed to a higher pO2, and where the carbide material has oxidized nearly to completion. Figure 10 shows lower magnification images of the oxidized regions characterized by porosity.   Figure 9: FIB lift out from the hot zone in (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar. [Left] Secondary electron image of the focus ion beam (FIB) cross section of part of the oxide scale. [Right] EDS maps showing elemental distribution of oxygen, titanium, tantalum, hafnium, zirconium and niobium. The white circle is added to help guide the eye to a partially oxidized grain.                    \\x0c', ' Figure 10: (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar. (a) Low magnification back-scattered electron image of the top half of the sample fracture cross-section; (b) Higher magnification image of the oxidized carbide cross section at the location indicated in (a); Oxygen EDS map of the region shown in (b).   3.3 Borides: After Oxidation Figure 11 shows low and high magnification BSE images of the oxide in plan view formed in the hot zone region of the diboride specimen, while Figure 12 shows secondary electron images of the oxidized region in cross-section. In contrast to the carbide, the diboride specimen exhibits a uniform scale in plan view. This scale is composed largely of Hf, Zr and Ti, with very little Ta and almost no Nb. Directly beneath this scale is a region with oxides and partially oxidized diborides. These regions exhibit porosity filled with a phase that has morphology consistent with that solidified from a melt (Figure 12), likely boria. Analysis of the FIB cross section (Figure 13) taken from the  oxidized (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 specimen shows that the surface oxide is dominated by Hf and Zr. Figure 14 shows a detail image of partially oxidized boride grains approximately 3 µm below the surface, depleted                   \\x0c', 'in Hf and Zr. The oxide around these grains is rich in Hf, Zr and Ti. A Ti rich region in the grain is also observed. The interface between the Ti-rich region and the center of the grain, which is depleted in Ti and enriched in Ta and Nb, is distinct, indicating the possibility of different boride phases. XRD results shown in Figure 15 and summarized in Table 2 confirm the preferential oxidation of group IV elements, and indicate the existence of a secondary group V rich boride.  Figure 11: (a) Back-scattered electron image showing the thin section of the (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 dogbone specimen oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar.; the orange curve shows the approximate size and location of the hot zone in plan view. (b) Higher magnification plan view image of the microstructure in the pyrometer sighting zone. (c) Semi-quantitative EDS results for the elemental composition of the oxide in plan view.    Figure 12: (Hf0.2ZrTi0.2Ta0.2Nb0.2)B2 oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar.. (a) shows a lower magnification fracture cross-section secondary electron image of the microstructure; (b) shows a higher magnification image of the region near the oxide/gas interface.                     \\x0c', ' Figure 13: (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar.. Backscattered electron image in cross section of the top of the oxide scale at the hot zone, shown in the top left along with EDS maps showing elemental distribution of boron, oxygen, titanium, tantalum, hafnium, zirconium and niobium.  The blue circular marker indicates an area also shown in Figure 14.   Figure 14: (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar. [Left] Secondary electron image of the focus ion beam (FIB) cross section of part of the oxide scale 3 µm below the surface in the hot zone, and EDS maps showing elemental distribution of boron oxygen, titanium, tantalum, hafnium, zirconium and niobium. The blue circular marker shown in Figure 13 is repeated again here for reference.                     \\x0c', '  Figure 15: X-ray diffraction spectra taken after oxidation for (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 oxidized for 5 mins at a nominal temperature of 1700°C in 1 mol% O2/bal Ar.  Table 2: Possible phase matches for the x-ray diffraction pattern (Figure 15) collected from (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 oxidized for five minutes at a nominal temperature of 1700°C in 1 mol% O2/bal Ar. PDF numbers or references are indicated in parentheses or brackets, respectively. Indicated Phase Crystal System Possible alternate phases/overlaps HfO2 (00-034-0104) Monoclinic (Hf, Zr)O2 [24,25] (Hf, Zr,Ti)O2  [23] ZrO2 (00-042-1164) Tetragonal (Hf, Zr)O2 [24,25] (Hf, Zr,Ti)O2 [23] (Zr, Ti)O2 (04-002-8273) Orthorhombic (Hf, Zr,Ti)O2 [23] (Hf, Zr)1-xTixO2 [23] Ta3B4 (04-003-3812) Orthorhombic (Nb,Ta)3B4 [26,28]                    \\x0c', '3.4 Thermodynamic Predictions Table 3 shows the results for the free energy minimization calculation for (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C and (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2 each in equilibrium with an equimolar solid solution containing HfO2, ZrO2, TiO2, TaO2.5, and NbO2.5, conducted at the experimental temperature, 1700°C. These thermodynamic predictions describe the local equilibrium at the carbide (or diboride) and oxide interface. Following the procedure described in detail in the companion paper [12], the calculation predicts the formation of an oxide primarily rich in the group IV elements (Hf, Zr and Ti), with Ta and Nb content in the oxide below experimentally determinable limits for both the carbide and the boride. Additionally, CO(g) is predicted to form with a fugacity of 0.99 atm for the carbide, while liquid boria is predicted to form for the diboride. The calculation for the high entropy carbide resulted in the prediction of a carbon deficient group V carbide, Nb8C7 (NbC0.87). Smith et al [27] reported the diffraction patterns for a range of rock-salt sub-stoichiometric niobium carbides above NbC0.71 and found the X-ray diffraction peaks to occur at the same or similar 2-theta values as stoichiometric NbC. Therefore, the peaks corresponding to “other carbide” shown in Figure 8 match multiple sub-stoichiometric niobium carbide reference patterns, hence the multiple possibilities indicated in Table 1.                          \\x0c', 'Table 3: Equilibrium calculations showing the final composition of the oxide scale and underlying substrate that results from an inputs of equimolar quinary alloy (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)C and  (Hf0.2Zr0.2Ti0.2Ta0.2Nb0.2)B2, respectively, at 1700°C in equilibrium with an equimolar (cation atom basis) oxide ideal solid solution containing HfO2, ZrO2, TiO2, Ta2O5, Nb2O5. † indicates separate phases predicted to form but that were not part of the substrate or oxide solid solution. Constituents Input (mol) Input (mol%) Carbide Output (mol %) Boride Output (mol %) Carbide/Boride Ideal Solution     Hf 1 20 0.015 0.025 Zr 1 20 0.022 0.180 Ti 1 20 14.1 15.5 Ta 1 20 44.7 42.2 Nb  1 20 41.1 42.2 Nb8C7 † CO (g) †(0.99 atm)  B2O3 (l) † Oxide Ideal Solution     HfO2  1 20 39.9 42.7 ZrO2  1 20 39.9 42.6 TiO2 1 20 20.2 8.50 Ti3O5 6.19 TaO2.5 1 20 8.22x10-4 9.89x10-3 NbO2.5 1 20 4.38x10-7 5.32x10-9 Equilibrium pO2 2.55x10-13 3.70x10-13  4.0 Discussion 4.1 Comparison of experimental results to calculated predictions The computational results shown in Table 3 predict the oxidation of group IV elements is thermodynamically favored and therefore preferentially oxidize, consistent with Part I of this work. As discussed in the companion publication [12], the thermodynamic prediction for the oxidized diborides                   \\x0c', 'are similar to that of the carbides, except that secondary, condensed phase, B2O3 (l), may also form, albeit with a high vapor pressure (~10-2 atm). The formation of a carbon-poor, Nb rich carbide phase is also predicted for the carbide system, in addition to the original rock-salt carbide phase. The formation of such a phase outside of the rock-salt carbide phase may be an artifact of the thermodynamic database used; prior studies or databases for carbides in the Nb-C system [27,29,30] do not have data for Nb8C7. A stoichiometry of 1:0.87 (Nb:C) in the phase diagram [29] at the test temperature indicates that the equilibrium phase is carbon-deficient NbC.  High temperature oxidation studies of niobium carbide have been observed to result in preferential oxidation of carbon [31], leading to metal-rich compositions; Smith et al. [27] have observed that for NbCx where x ≤ 0.63, the niobium carbide adopts a trigonal structure.   The experimental observations are in good agreement with the calculated predictions for the oxide composition formed on the carbide and the diboride at the substrate/oxide interface: the group IV elements preferentially oxidize over the group V elements, while the substrate is enriched in group V elements. Within these groups, the extent of oxidation of the different elements also corresponds to their oxide thermodynamic favorability. Hf and Zr, which form the most stable oxides, oxidize preferentially, followed by Ti. Of the group V elements, Ta2O5 is more stable than Nb2O5, and thus Ta is expected to preferentially oxidize over Nb.  The combination of intergranular oxidation and preferential oxidation would result in an oxide scale with an evolving composition until the grain is completely consumed. This explains the other phases formed after oxidation of the carbide and detected via XRD (Figure 8, Table 1) on the specimen surface where the oxidation of grains has gone to completion or near completion. The Ta content in the oxide formed on the carbide is higher than both the calculated predictions and the experimental observations for the oxidized boride. Oxidized regions with higher Ta content than predicted by the                   \\x0c', 'thermodynamic calculations were observed in regions farther away from the carbide/oxide interface. As the group IV elements are depleted, the substrate becomes group V rich; further oxidation of the grains therefore results in the incorporation of Ta in the oxides. Comparison of the oxidation of the carbides and diborides indicate that this occurred for the carbide for the oxidation conditions studied in this work, but not the diboride. The porous, intergranular oxides formed on the carbide (Figure 9, Figure 10) is attributed to the formation and release of CO(g) [32,33], which is the dominant gaseous oxide product formed upon oxidation of the carbide (Table 3). In comparison, the formation and retention of liquid boria attending the oxidation of diborides at this temperature may seal pores and defects in the oxide and result in a more effective barrier to oxidant ingress [2,34,35] compared to the gas-filled pores in the carbide case. Such a pore-filling phase was observed (Figure 12) in the oxidized diboride beneath the group IV-rich scale. The composition of this phase was indeterminate via EDS analysis due to poor counts. The morphology of this phase is consistent with solidification from a melt, and is inferred to be liquid boria, which is known to be retained in the scale of oxidized diborides up to 1800°C [35]. Other possible oxide or boride phases as determined from EDS and XRD (Figure 15, Table 2) do not have melting temperatures at or below 1700°C.  The formation of a porous scale upon oxidation of the carbide also results in higher oxygen partial pressures at the reaction interface in the carbide, which allows for the further oxidation of Ta over Nb (Ta2O5 is more stable than Nb2O5). Figure 7 and the EDS maps in Figure 11 show that Ta is present in some of the oxygen rich regions, but almost no Nb. XRD matches for Nb-rich carbides are consistent with this interpretation and calculated predictions. In contrast, the dense surface scale formed on the diboride and filled pores likely reduces the oxygen partial pressure at the reaction interface. This explains not only the predicted low tantalum content in the oxide formed on the diborides, but also the titanium retained in the boride phase in Figures 11 and 12. Hf and Zr oxides are significantly more                   \\x0c', 'thermodynamically stable than TiO2 (see Ellingham diagram in companion paper [12]); Hf and Zr oxidize preferentially from the boride, leading to Ti “lagging” the other group IV elements in extent of oxidation. The depletion zone in the Ti map (Figure 15) indicates that some Ti is still retained in the diboride. The Ta3B4 boride phase indicated by the XRD results is consistent with a tantalum rich composition, an analog to Nb rich substoichiometric carbide in the high entropy carbide case.   4.2 Complex oxide formation The preceding work and that outlined in the companion paper are predicated on the assumption that the carbide or boride is in equilibrium with a solid solution oxide. The lack of available thermodynamic data for complex or ordered compounds limited consideration of these phases. A review of the relevant phase diagrams indicate that only Hfand Zroxides exhibit complete solubility. Other binary and/or ternary oxide systems under consideration exhibit only partial solid solubility and complex oxide formation. Further, the XRD patterns for the oxidized samples indicate that oxide compounds such as Zr6Nb2O17 and (Hf,Zr)TiO4 formed. Table 4 summarizes known complex oxides which are likely to form based on a review of the available phase diagrams. As presented in section 3.0, the actual oxides formed on the high entropy ceramics are complex variations of the compounds shown in Table 4, due to the possibility of ionic substitutions. For example, Hf and Zr are known to form substitutional solid solutions in alloys and oxides, and in their computational work, Hautier et al [26]found that Ta and Nb have a high likelihood to substitute for each other. In fact, an analogous phase to Zr6Nb2O17 exists with Hf and Ta: Hf6Ta2O17 [22,36].   It has been posited by Butler et al  [9] and Gorr et al [37] that the formation of complex oxide phases may promote improved oxidation resistance, and therefore be an advantage of high entropy materials relative to conventional alloys and ceramics. Consideration of the complex oxide phases was                   \\x0c', 'outside the scope of this work, largely due to the lack of thermodynamic data. However, it can be seen from the experimental results that despite the formation of complex oxides, the elemental composition of the assemblage of oxides still followed the predicted trend, wherein the group IV elements oxidized preferentially and dominated the composition of the thermally grown oxide.  Table 4: Known complex oxides that could form in the material systems under study, based on a review of available phase diagrams.   Hf Ti Zr Nb Ta Hf           Ti  HfTiO4 [38]          Zr   ZrTiO4  [23,39,40]       Nb Hf6Nb2O17  [41] Nb2TiO7  [42] Nb10Ti2O29 [42] Nb6Ti2O19   [42] TiNb6O17 [43] Zr6Nb2O17 [43]         Ta Hf6Ta2O17  [36] TiTa2O7 [21] ZrTa6O17 [44] Zr6Ta2O19 [44] Nb4Ta2O15 [45]   V     ZrV2O7 [46] VNb9O25  [47] VTa9O25 [48,49] Mo           W HfW2O8 [50]   ZrW2O8  [50]      4.3 Implications for use of HE-UHTCs in oxidizing environments Oxides formed from group IV elements (Hf, Zr, Ti) have the highest melting temperatures (Tm) and are most thermodynamically favored among the refractory elements. The predicted and observed preferential oxidation of these elements demonstrate that an HE-UHTC can be designed to form an oxide scale containing these elements.  Further, the retention of elements with low Tm oxides as carbides or borides ensures a material with a high thermal stability overall, as Nb and Ta carbides and borides have melting temperatures greater than 2000°C [51].  The approach outlined in Part I and this work can                   \\x0c', 'further be used to optimize such a composition as a multi-principal component material that is not equimolar, but promotes a targeted oxide and substrate composition.  On the other hand, the preferential depletion of several elements from the substrate can destabilize the underlying substrate and promote the formation of other phases. This destabilization can result in the loss of properties that might be gained from the high entropy design concept.   5.0 Conclusions  The objective of this work was to test the predictions of the analytical and computational thermodynamic approach outlined in a companion publication. This methodology was applied to the prediction of substrate and oxide composition after oxidation of a high entropy carbide and boride containing Hf, Zr, Ti, Ta and Nb in equimolar amounts. The experimentally determined oxide and substrate compositions were found to be in good agreement with the thermodynamic predictions, which describe well the local equilibrium at the reaction interface for both the carbide and the diboride case. In general, the oxidation of high entropy materials studied here occurred by preferential oxidation of each component according to the relative thermodynamic favorabilities of their respective oxides. Group IV oxides, being most favorable, formed preferentially. The experimental results further elucidated that within each group, preferential oxidation can also occur according to relative thermodynamic stability: Hf and Zr preferentially oxidize over Ti as shown in the boride case, while Ta preferentially oxidizes over Nb as shown in the carbide case. The substrate composition of group V metal-rich carbides and borides, which are or have the potential to form phases with different crystal structures from the original ceramic, was accurately predicted. Whereas complex oxides were observed experimentally, this is not a significant limitation to the prediction of preferential oxidation, as the formation of these complex oxides was dominated by predicted oxide constituents.                    \\x0c', ' This work shows that the approach outlined in Part I provides insight into the thermodynamics driving the oxidation behavior of high entropy materials and can be used as a tool for the design of high entropy materials for oxidation resistance.   Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.   6.0 Acknowledgements  This work is supported by the U.S. Office of Naval Research MURI program (grant no. N00014-151-2863) and the Virginia Space Grant Consortium Graduate Research Fellowship. Bulk carbide samples were provided by the Vecchio research group at the University of California San Diego. The authors would also like to thank Professors Bill Fahrenholtz (Missouri University of Science and Technology), Bill Soffa, Bi-Cheng Zhou (University of Virginia) and Christina Rost (James Madison University) for helpful discussions, and Dr. Helge Heinrich (University of Virginia) for help with the FIB lift-outs. Characterization was conducted at the Nanoscale Materials Characterization Facility at University of Virginia.                      \\x0c', 'References [1] D. 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Coughanour, Phase equilibrium relations in the systems titania-niobia and zirconia-niobia, Journal of Research of the National Bureau of Standards. 55 (1955) 209. https://doi.org/10.6028/jres.055.023. [44] R. Roth, J. Waring, Effect of oxide additions on the polymorphism of tantalum pentoxide III. Stabilization of the low temperature structure type, J. Res. National Bureau of Standards-A. Physics and Chemistry. (1970) 485-493. [45] F. Holtzberg, A. Reisman, Sub-Solidus Equilibria in the System Nb2O5 -Ta2O5, The Journal of Physical Chemistry. 65 (1961) 1192-1196. https://doi.org/10.1021/j100825a024.  [46] V. Cirilli, A. Burdese, C. Brisi, Atti Accad. Sci. Torino, Cl, Sci. Fis., Mat. Nat. 95 (1961) 197-228. [47] J. Waring, R. Roth, Phase equilibria in the system vanadium oxide-niobium oxide, J. Res. Natl. Bur. Std. 69 (1965). [48] R.S. Roth, J.L. Waring, W.S. Brower, Effect of oxide additions on the polymorphism of tantalum pentoxide. II. “„Stabilization‟” of the high temperature structure type, Journal of Research of the                   \\x0c', 'National Bureau of Standards Section A: Physics and Chemistry. 74A (1970) 477. https://doi.org/10.6028/jres.074A.037. [49] H. Schadow, H. Oppermann, B. Wehner, Investigations on the Quasi‐binary System V2O5-Ta2O5, Crystal Research and Technology. 27 (1992) 691-695. [50] L.L. Chang, M. Scroger, B. Phillips, Condensed phase relations in the systems ZrO2‐WO2‐WO3 and HfO2‐WO2‐WO3, Journal of the American Ceramic Society. 50 (1967) 211-215.  [51] E. Rudy, Ternary phase equilibria in transition metalboron-carbon-silicon systems. Part 5Compendium of phase diagram data Summary report, 1 Jan. 196430 Apr. 1969 (Phase diagrams of binary transition metal systems and binary and ternary systems of refractory metal alloys), (1969).                   \\x0c', 'Graphical Abstract                    \\x0c']"
},{
  "_id": 75,
  "PDF": "Fabrication and evaluation on thermal stability of hafnium diboride matrix composite at severe oxidation condition.pdf",
  "Text": "['Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  Contents lists available at ScienceDirect  Int.  Journal of Refractory Metals & Hard Materials  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / I J R M H M  Fabrication and evaluation on thermal stability of hafnium diboride matrix composite at severe oxidation condition  Ling Weng a,b,*, Xinghong Zhang b, Wenbo Han b,  Jiecai Han b  a The School of Material Science and Engineering, Harbin University of Science and Technology, Harbin 150040, People’s Republic of China b Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, People’s Republic of China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 27 October 2008 Accepted 3 December 2008  Keywords:  Ceramics Oxidation Microstructure  1. Introduction  Two hafnium diboride based composites, respectively, containing 20 vol%SiC and 20 vol%SiC-10 vol%AlN, were hot-pressed. Microstructures and mechanical properties were investigated and the thermal stability over 2000 °C was evaluated by oxyacetylene torch. Results indicated that the addition of AlN greatly improved the powder sinterability by eliminating the oxygen contamination. The mechanical properties for HfB2-SiC-AlN composite, especially the flexural strength were enhanced remarkably through the improvement in density and the formation of AlN-SiC solid solution. Oxyacetylene ablation results indicated that the addition of AlN also improved the thermal stability of composite under severe oxidation conditions. The ablation behavior was investigated and the ablation mechanism was discussed. Ó 2008 Elsevier Ltd. All rights reserved.  IVA group transition metal diborides such as ZrB2 and HfB2, which were called the ultra high temperature ceramics (UHTCs), have been indicated as promising candidate materials for use in some aerospace applications, primarily for melting temperatures greater than 3200 °C. Other favorable characteristics include high elastic modulus, high thermal conductivity, retained strength at elevated temperature, relatively good thermal shock resistance and modest thermal expansion [1]. In view of high temperature applications, an understanding of thermal stability at high temperature is a key factor for the optimization of the material performances. Refractory nitride (i.e. AlN, Si3N4, etc.) had been reported as effectively sintering additives in fabrication the refractory metal diboride ceramics with good performances [2-4]. These refractory nitrides could effectively reduce the occurrence of secondary grain boundary phases than metal additives, resulted some greatly improvement in the sinterability, microstructures and properties of UHTCs. However, the mechanism of refractory nitrides as sintering aids was not yet understood well for researchers and the effects of refractory nitrides on the properties of UHTCs especially the ablation resistance need to be investigated deeply. A number of researchers had worked on the oxidation resistance of HfB2-based and ZrB2-based composite under the condi * Corresponding author. Address: The School of Material Science and Engineering, Harbin University of Science and Technology, Harbin 150040, People’s Republic of China. Tel./fax: +86 451 86392533. E-mail address: wengling79@163.com (L. Weng).  0263-4368/$ see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2008.12.001  tions on static air and ambient pressure [5-7]. However, these oxidation conditions could not effectively simulate the truly flight conditions. In recent years, some large size arc jet tests [8] have been performed in evaluating the thermal behavior of monolithic HfB2 or ZrB2 ceramics at actual flight simulations, but the enormous experimental cost, complicated experimental process and the difficulty in fabricating the large size samples hindered the widely application of these tests. So a simple and economical method should be developed to effectively evaluate the thermal stability of Zror Hfbased ceramics at severe oxidation conditions. Oxyacetylene torch facility could provide a low cost, simplicity and rather good simulated conditions on evaluating the ablation resistance behavior of UHTCs, including high temperature (>2000 °C) and severe oxidation environment. It could help researchers to obtain a better understanding on the oxidation and thermal performance of ZrB2 or HfB2 based composites. In this work, two HfB2-SiC composites, doped and undoped with AlN, were fabricated by hot-pressing sinter. The microstructure and mechanical properties of these two composites were investigated. The thermal ablation resistances were evaluated by oxyacetylene torch test.  2. Experimental  Commercially available HfB2 powder (size 5-7 lm, 97.1% purity, oxygen content 2.28 wt.%, other trace contamination content 0.62 wt.%. Northwest Institute for Non-Ferrous Metal Research, China), SiC powder (purity > 99.5%, size 2-4 lm, Xuzhou Hongwu Nanometer Materials Co. Ltd., China) and AlN powder  \\x0c', '712  L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  Table 1 Experimental condition for oxyacetylene ablation.  O2 gas pressure (kPa) O2 gas flux (L/s) C2H2 gas pressure (kPa) C2H2 gas flux (L/s) Diameter of nozzle (mm) Distance from sample surface to nozzle (mm)  0.45 1.02 0.1 0.36 2 10  (purity > 98%, particle size <100 nm, Hefei Kiln Nanometer Technology Development Co. Ltd., China) were used. Two mixtures made from commercially available (vol%):  powders  (1) HfB2-20SiC (abbreviated as HS) (2) HfB2-20SiC-18AlN (abbreviated as HSA)  were milled for 24 h in polyethylene jars with pure ethanol as solvent and WC balls as milling balls, dried with a rotating evaporator and sieved by 200 mesh sieve. The powder mixtures were uniaxially hot-pressed in Ar atmosphere. Sintered temperature/ dwell times/applied pressures were: 2200 °C/60 min/30 MPa for HS and 1800 °C/30 min/30 MPa for HSA, respectively. The final density of composites was measured by Archimedes water-immersion method, the theoretical density was estimated by the rule of mixture. Flexural strength and Fracture toughness  Fig. 2. SEM-EDS examination of was highlighted.  the polished section of HSA. A Si-Al-C-N phase  (KIC) were tested by three points bending and single-edge notched beam at Electronic Universal Testing Machine (Instron 5500). At least five specimens were tested for each experimental condition. The thermal stability was tested by oxyacetylene ablation test. Table 1 listed the specific experimental condition. During the test, a specimen 15 mm in diameter and 10 mm in height was exposed to the flame. Ablation time was 300 s. The distance between the  Fig. 1. SEs-SEM micrographs from polished surfaces of HS (a) and HSA (b). Arrows indicated the micro-cracks between grains.  Fig. 3. The fractural mode of the composite HS (a) and HSA (b).  \\x0c', 'L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  713  Table 2 Mechanical properties of the composites HS and HSA.  Composite  Relative density (%)  Hv1.0 (GPa)  KIC (MPa m1/2)  HS HSA  93.8 99.6  18.3 ± 0.5 18.5 ± 0.4  4.6 ± 0.6 5.0 ± 0.3  r (MPa)  353 ± 25 612 ± 30  C  º  /  nozzle tip of the oxyacetylene gun and the surface of the specimen was 10 mm. To ensure ablation in just one direction, specimen was held in a concave graphite anvil and only one surface was exposed to the ablation environment. The flow rates of oxygen and acetylene were 1.02 and 0.36 L/s, respectively. The inner diameter of the nozzle tip of the oxidation gun was 2 mm. Surface temperature of sample was measured using an optical pyrometer and recorded every 10 s. The mass equations:  calculated  rate was  following  ablation  the  by  e  r  u  t  a  r  e  p  m  e  T  2500  2000  1500  1000  500  0  HS HSA  0  50  100  150  Time /s  200  250  300  md ¼ jm2 \\x00 m1 j t  ð1Þ  Fig. 5. The temperature schedule of the tests.  where md is the mass ablation rate, m1 and m2 represent the mass of samples before and after the oxyacetylene ablation test, t represents the ablation time. Surfaces and polished cross-sections were analyzed by scanning electron microscope (SEM, Philips, 4700) and energy dispersive microanalysis.  3. Results and discussion  3.1. Microstructure and mechanical properties  The typical grained structure of the hot-pressed ceramics was given in Fig. 1. SEM image of the polished surface of composite HS showed regularly shaped HfB2 grains (10 lm maximum in size) and SiC particles. Some micro-cracks were observed at the grain boundary between HfB2 and SiC grains, as arrows pointed. The occurrence of microcracking events could be attributed to the emergence of high residual stresses which come up at the thermal expansion mismatch between HfB2 and SiC. Composite HSA was characterized by finer grains (about 5 lm) and more compact structure than composite HS. The SEM-EDS examination of the polished section identified the two basic phase (HfB2 and SiC), it also highlighted a more minority Si-Al-C-N phase (Fig. 2), this new phase was possibly a solid solution phase of AlN-SiC. This fine grained microstructure might be attributed to a lower temperature and shorter sintering time, for which the abnormal growth of grains was prevented. Moreover, the occurrence of AlN-SiC solid solution also accelerated the sintering process at rather low sintering temperature through the liquid phase transmission mechanism. Fig. 3 showed the different fracture mode of two  composites, in which one could clearly observed that composite HS mainly fracture at transgranular mode while composite HSA at inter/trans-granular fracture mode. Some thermal-mechanical properties of two composites were listed in Table 2. The relative density of composite HS was only 93.8%, which was greatly lower than HSA (99.6% at r.d). Limited three point flexural tests and fracture toughness tests at room temperature were carried out. For composite HS, the flexural strength and fracture toughness were 353 MPa and 4.6 MPa m1/2 respectively, which were much lower than the corresponding values for composite HSA (612 MPa and 5.0 MPa m1/2). This was possibly due to the higher relative density and finer grains of HSA. The improvement in mechanical properties of composite HSA when comparing to the composite HS, could be attributed to two possible reasons: one was the densification effect of AlN. Monteverde [2] investigated the densification of diborides was greatly hindered by some oxygen contamination, such as B2O3, HfO2 or their mixture, due to the reduction of boron activity. The decrease in boron activity was accompanied by a reduction of the densification rate which, in turn, promoted grain coarsening by evaporationcondensation kinetics. Presumably, AlN enabled densification in HfB2 by facilitating the removal of HfO2 and B2O3 at temperatures low enough to prevent significant coarsening of the HfB2 grains before densification. The reactions were listed below [2]:  2AlN þ B2O3 ! Al2O3 þ 2BN  8AlN þ 6HfO2 ! þ6HfN þ N2  ð2Þ  ð3Þ  Fig. 4. The macrographs of HS (a) and HSA (b) after 300 s oxyacetylene torch test.    \\x0c', '714  L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  Fig. 6. The SEM micrographs of the oxidized surface for HS and HSA: (a) low magnitude SEM for HS; (b) high magnitude SEM for HS; (c) low magnitude SEM for HSA; (d) high magnitude SEM for HSA;  Fig. 7. Cross-section SEM microstructure analysis of HS and HSA composites after 300 s of the oxyacetylene test: (a) HS; (b) HSA. The whole oxidation zones were mainly divided into three-layers (from top to buttom): outermost oxide layer (1), oxide sub-layer (2) and unreacted bulk (4).  \\x0c', 'L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  715  Fig. 8. SEM images of oxide sub-layer of the two composite. (a) low magnitude of HS oxide sub-layer, (b) high magnitude of HS oxide sub-layer, (c) low magnitude of HSA oxide sub-layer, (d) high magnitude of HSA oxide sub-layer, EDS analysis indicated the Hf,-Al-O phase.  Another possible reason was the formation of AlN-SiC solid solution, as showed in Fig. 2. According to references [9-12] and the tentative phase diagram of AlN-SiC system [13], silicon carbide and aluminium nitride could form a series of solid solutions in a very wide range of composition and it could be sintered, preferably by hot pressing in the temperature between 1700 °C and 2300 °C. This solid solution had a better combined on thermal-mechanical properties than monophase AlN or SiC ceramics. So the reinforced effect of AlN-SiC solid solution in HfB2 bulk was better than SiC particle along.  3.2. Thermal stability at severe oxidation condition  Fig. 4 showed the macrographs of HS and HSA after exposure to the oxyacetylene torch flame for 300 s. Fig. 5 showed the increasing temperature schedule during the oxyacetylene tests. It should be noted that the surface temperature of HSA was slightly lower  Table 3 Possible oxidation equations and their Gibbs freen energy in the system.  No.  (4) (5) (6) (7) (8) (9)  Formulas  HfB2 + 5/2O2 = HfO2 + B2O3 SiC(s) + 2O2 (g) = SiO2 (s) + CO2 (g) SiC(s) + 3/2O2 (g) = SiO2 (s) + CO(g) 2SiO2 (s) + SiC(g) = 3SiO(g) + CO(g) B2O3 (l) + H2O(g) = 2HBO2 (g) B2 O3 (l) + 3H2O(g) = 2H3BO3 (g)  Dr G0  m (KJmol\\x001) m ¼ \\x002047:65 \\x00 0:44297T m ¼ \\x001225:49 \\x00 0:16747T m ¼ \\x00942:52 \\x00 0:08088T m ¼ 1471:82 \\x00 0:72276T m ¼ \\x00111:56 \\x00 0:09922T m ¼ \\x00211:42 \\x00 0:46716T  Dr G0 Dr G0 Dr G0 Dr G0 Dr G0 Dr G0  than that of HS, which was possibly caused by the higher thermal conductivity for the former, due to the addition of AlN. For HS, the flow pattern of the oxidation surface showed a mechanical scour resulted from the blowing by the oxyacetylene flame, corresponding to a large mass erosion rate (0.0267 g/s). The more compact and homogeneous oxide scale was formed on the HSA composite, no cracks were observed during oxidation or after being cooled down from an elevated temperature, the composite showed an obviously lower mass erosion rate (0.01477 g/s). Further observation by SEM (Fig. 6) showed the very different microstructures of the oxidized surface for HS and HSA. Extensive bubble formation composed of 2-3 lm particles was observed on HS sample surface. EDS analysis indicated that these particles were HfO2 grains, which was the crystallization of molten HfO2. Few silicon glassy phase was observed on the oxidized surface of HS. It might be attributed to the strong airflow acted on the sample surface which could greatly destroyed the formation of SiO2 glassy phase then decreased the oxidation property of composite. The micrographs in Fig. 6c and d indicated the representative microstructural modifications of the HSA surface after acetylene ablation test over 2000 °C. The surface was covered by a thick glassy layer, bubbles with different sizes were observed in this glassy layer. It should be noted that this glassy layer was mainly consist of Si, Al and O, which was not reported in previous studies. The formation of Si-Al-O glassy layer was most likely due to the oxidation of SiC and AlN or its solid solution on the sample surface. This liquid solution of Al2O3 in SiO2 was in agreement with the SiO2-Al2O3 phase diagram [14].  \\x0c', '716  L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  Oxidation atmosphere  O2  H2O  Unaffected bulk   CO, SiO, B2O3,   HBO2, H3BO3  O2  H2O   (1)   (2)   Pore  Unaffected bulk   Oxidation atmosphere   Oxidation atmosphere   O2   H2O   O2  H2O   (1)   (2)   (1)   (2)   Pores filled by the   Hf-Al-O glass phase   (3)   Unaffected bulk   (3)   Unaffected bulk   HS   HSA   Fig. 9. Schematic graphs of oxidation model for HS and HSA.  The cross-sections of composite HS and HSA (Fig. 7) after 300 s oxyacetylene ablated were examined to determine changes that occurred in the microstructure near the ablation surface. Both of the cross section morphologies of two composites were composed of a three-layer structure from top to bottom: total oxidation layer, oxide sub-layer and unaffected bulk. Fig. 8 showed the oxide sublayer of the two composites. From Fig. 8a and b, one could clearly observe that the oxide sub-layer of composite HS was a porous structure, which was attributed to the oxidation of SiC particles. When further investigated, some interesting differences were founded on oxide sub-layer between composite HS and HSA, in which a continuous glassy phase was observed in the latter (Fig. 8c and d). Energy spectrum analysis indicated that this glassy phase was mainly composed of three elements: Hf, Al and O. It should be noted that EDS spectra also showed a carbon peak around 0.18 keV, which might be attributed to the active oxidation of SiC. Monteverde [15] had reported that it was as much as plausible that conditions of high temperature and low oxygen partial pressure in the oxidized sub-layer were met for the active oxidation of SiC [16-18]. From Fig. 7b, it could be also observed that this Hf-Al-O glassy phase tended to decrease from top to bottom. For the HfB2-SiC based UHTCs, the expected main reactions describing the oxidation process involve the oxidation of ceramic with oxygen and the combination of H2O vapor with some oxidation products, which were listed in Table 3. (The thermodynamic data obtained from Ref. [19].) In terms of the principle of minimum free enthalpy, the reaction would occur spontaneously when the Gibbs free energy was negative. As showed in Table 3, all the exothermal reactions could occur spontaneously due to Dr G0 m < 0: After AlN addition, the following reaction possibly proceed:  4AlNðsÞ þ 7O2 ðgÞ ¼ 2Al2O3 ðsÞ þ 4NOk2 ðgÞ  ð10Þ  The addition of SiC resulted in the formation of many low melting point liquids, low viscosity glasses and gas phases, as shown in above reactions. The fluid phases containing Si readily peeled out  under the combustion gas impact. Hence the ablation properties of the HS would be deteriorated with the SiC addition. The effect of AlN on the ablation behavior of the HSA could be attributed to the formation of high melting point Al2O3 which had a high vapor pressure and low volatility at high temperature, and it could in situ block the attack by the combustion gas. Fig. 9 indicated the microstructural development of the oxide layers of two composites during the ablation process. The improved oxidation resistance mechanism of the composite with added AlN as sintering aid (HSA) could be depicted as follow: During the ablation process, combined with some oxidation reactions such as the oxidation of HfB2, active or passive oxidation of SiC and the evolution of volatile products (B2O3, CO, SiO), a porous structure of oxide layers was occurrence. For composite HSA, the occurrence of HfO2-Al2O3 glassy phase could spread over the oxidation surface and prevented the inward penetration of O2, due to its low viscosity and good fluidity, then provided a better protection on the oxidation layer than HfO2 alone. Furthermore, the glassy phase could fill in the voids in the HfO2 framework and endow the oxidation layer with a moderate strength, preventing it from being broken off.  4. Conclusion  HfB2-SiC composites with and without AlN addition were fabricated by hot-pressed sintering. The microstructures and mechanical properties of two composites were investigated. The thermal stability was examined by oxyacetylene ablation tests. Results indicated that the thermal-mechanical properties of HfB2-SiC- AlN composite were much better than that of HfB2-SiC composite. These increases on mechanical properties of the composite with AlN addition could be attributed to the great improved in relative density and reduced in grain size. The composite HSA had a better ablation resistance than composite HS, as evaluated by comparison of the fractional mass loss. SEM observations of the cross-sections showed a three-layers structure. A clue of redistribution of Hf-Al-  \\x0c', 'L. Weng et al. / Int.  Journal of Refractory Metals & Hard Materials 27 (2009) 711-717  717  O glassy phase which was observed in the oxide sub-layer of composite HSA was considered to have a positive effect on the ablation resistance of composite HSA.  Acknowledgement  This work was supported by the National Natural Science Foundation of China (90505015 and 50602010), the Research Fund for the Doctoral Program of Higher Education (20060213031) and the Program for New Century Excellent Talents in University.  References  [1] Monteverde F, Savino R. Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions. J Eur Ceram Soc 2007;27: 4797-805. [2] Monteverde F, Bellosi A. Beneficial effects of AlN as sintering aid microstructure and mechanical properties of hot-pressed ZrB2. Adv Mater 2003;5:508-12. [3] Monteverde F, Bellosi A. Efficacy of HfN as sintering aid in the manufacture of ultrahigh temperature metal diborides matrix ceramics. J Mater Res 2004;19: 3576-81. [4] Monteverde F, Bellosi A. Effect of behavior and microstructure 2002;46:223-8. [5] Brach M, Sciti D, composite in the 1771-80.  the addition of silicon nitride on sintering of zirconium diboride. Scripta Mater  Bellosi A. Short-term oxidation AlN-SiC-ZrB2. J Eur Ceram  a ternary 2005;25:  Balbo A, system  on Eng  of Soc  [11]  [6] Tabaru T, Shobu K, Sakamoto M, Hirai H, Hanada S. Oxidation behavior of Mo(Si0.6,Al0.4)2/HfB2 composites as aluminium reservoir materials for protective Al2O3 formation. Scripta Material 2003;49:767-72. [7] Levine SR. Tantalum addition to zirconium diboride for improved oxidation resistance. 2003; NASA/TM-2003-212483. [8] Gasch M, Ellerby D, Irby E, Beckman S, Gusman M, Johnson S. Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra-high temperature ceramics. J Mater Sci 2004;39:5925-37. [9] Rafaniello W, Cho K, Virkar AV. Fabrication and characteristics of SiC-AlN alloys. J Mater Sci 1981;16:3479-88. [10] Raffenielo W, Plichta MR, Virkar AV. Investigation of phase stability in the system SiC-AlN. J Am Ceram Soc 1983;66:272-6. Jou ZC, Kuo SY, Virker AV. Elevated temperature creep of silicon carbide- aluminium nitride ceramics: role of grain size. J Am Ceram Soc 1986;69:279-81. [12] Ruh R, Zangvil A. Composition and properties of hot pressed SiC-AlN solid solutions. J Am Ceram Soc 1982;65:260-5. [13] Zangvil A, Ruh R. Phase relationships in the silicon carbide aluminium nitride system. J Am Ceram Soc 1988;71:884-90. [14] Zhou Y. Ceramic materials. 3rd ed. Beijing: Science Press; 2004. [15] Monteverde F, Scatteia L. Resistance to thermal shock and to oxidation of metal diborides-SiC ceramics for aerospace application. J Am Ceram Soc 2007;00:1-9. [16] Chamberlain A, Fahrenholtz WG, Hilmas G, Ellerby D. Oxidation of ZrB2-SiC ceramics under atmospheric and re-entry conditions. Refrac Appl Transact 2005;2:1-8. [17] Fahrenholtz WG. Thermodynamic analysis of ZrB2-SiC oxidation: Formation of a SiC depleted region. J Am Ceram Soc 2007;90:143-8. [18] Heuer AH, Lou VLK. Volatility diagrams of silica, silicon nitride and silicon carbide. J Am Ceram Soc 1990;73:2789-803. [19] Ye DL, Hu JH. Handbook of thermodynamic Metallurgical Industry Press; 2002.  inorganics.  Beijing:  data  of  \\x0c']"
},{
  "_id": 76,
  "PDF": "Fabrication and properties of ZrB2-SiC-BN machinable ceramics.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 1035-1042  Fabrication and properties of ZrB2-SiC-BN machinable ceramics  Haitang Wu a,b , Weigang Zhang a,∗  a State Key Laboratory of Multi-Phase Complex Systems, Institute of Process Engineering, Chinese Academy of Sciences, Beijing 100190, PR China b Graduate University of Chinese Academy of Sciences, Beijing 100049, PR China  Received 16 April 2009; received in revised form 5 September 2009; accepted 22 September 2009  Available online 29 October 2009  Abstract     ZrB2 -SiC-BN ceramics were fabricated by hot-pressing under argon at 1800 C and 23 MPa pressure. The microstructure, mechanical and oxidation resistance properties of the composite were investigated. The flexural strength and fracture toughness of ZrB2 -SiC-BN (40 vol%ZrB2 -25 vol%SiC-35 vol%BN) composite were 378 MPa and 4.1 MPa m1/2 , respectively. The former increased by 34% and the latter decreased by 15% compared to those of the conventional ZrB2 -SiC (80 vol%ZrB2 -20 vol%SiC). Noticeably, the hardness decreased tremendously by about 67% and the machinability improved noticeably compared to the relative property of the ZrB2 -SiC ceramic. The anisothermal and isothermal oxidation behaviors of ZrB2 -SiC-BN composites from 1100 to 1500 C in air atmosphere showed that the weight gain of the 80 vol%ZrB2 -20 vol%SiC and 43.1 vol%ZrB2 -26.9 vol%SiC-30 vol%BN composites after oxidation at 1500 C for 5 h were 0.0714 and 0.0268 g/cm2 , respectively, which indicates that the addition of the BN enhances oxidation resistance of ZrB2 -SiC composite. The improved oxidation resistance is attributed to the formation of ample liquid borosilicate film below 1300 C and a compact film of zirconium silicate above 1300 C. The formed borosilicate and zirconium silicate on the surface of ZrB2 -SiC-BN ceramics act as an effective barriers for further diffusion of oxygen into the fresh interface of ZrB2 -SiC-BN. © 2009 Elsevier Ltd. All rights reserved.              Keywords: Zirconium diboride; Boron nitride; Oxidation resistance; Mechanical properties; Machinability  1.  Introduction  Ultra-high temperature ceramics (UHTCs) including refractory diborides and carbides, such as ZrB2 , HfB2 , ZrC, HfC and TaC, are considered as the most promising materials for the application in critical thermal protection systems and other components of future hypersonic aircraft or re-entry vehicles. Compared to the single-phase monolithic UHTC, ZrB2-SiC composite is of particular interest because of its striking property combination of high melting point, resistance to ablation/oxidation at high temperatures, high electrical and thermal conductivity and thermal-shock resistance, which makes it an attractive potential candidate for aerospace applications.1-5 Machining is an inevitable requirement for flexible use of advanced ceramics, especially when the complex and preci ∗  Corresponding author at: State Key Laboratory of Multi-Phase Complex Sys tems, Institute of Process Engineering, Chinese Academy of Sciences, Beijing  100190, PR China. Tel.: +86 10 62520135; fax: +86 10 62520135.  E-mail address: wgzhang@home.ipe.ac.cn (W. Zhang).  0955-2219/$ - see front matter © 2009 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2009.09.022  sion components of structural ceramics are involved. The use of diamond tools and some special processing technologies such as laser machining and ultrasonic machining are often inefficient and costly (machining cost usually accounts for 70-90% of the total cost) though those processes make some hard ceramics materials machinable.6 Electrical discharge machining is another promising technology to machine ceramic components of complex shape with high-dimensional accuracy and low surface roughness. Except for the inefficiency, electrical discharge machining requires a material resistance and can only machine components of small size.7 Comparatively, traditional mechanical machining is of both cost-effective and time-efficient. However, the extremely high strength and hardness of ZrB2-SiC due to the coexistence of strong covalent and metallic bond make mechanical machining very difficult or even impossible. This prevents the material from wide application. In recent years, attempts have been made to improve the machinability of ceramic materials by introducing in the matrix weak interfaces material, such as mica, h-BN, graphite, pores, rare-earth phosphates and Ti3SiC2 analogous compounds, to facilitate crack deflection during machining.8 Among those materials, h \\x0c', '1036  H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  BN, which exhibits high thermal conductivity and high melting point, is regarded as a suitable and effective interface material since the cleavage plane of h-BN facilitates crack propagation and decreases the cutting resistance during machining.9-12 Besides these, the similar crystal structures of BN and ZrB2 ensure the good chemical compatibility between them. Therefore, ZrB2-SiC-BN system may be a good candidate material for high temperature ceramics with excellent machinability and mechanical properties. In this paper, the fabrication of a machinable ZrB2-SiC-BN composite with excellent mechanics properties and oxidation resistance properties was reported. Influences of BN content on the hardness, fracture toughness, flexural strength and oxidation resistance property of the composite were investigated. The oxidation resistance properties tested at from 1100 to 1500 C were reported here. As the operative temperature for UHTC materials is in excess of 2000 C, composition performing better than others at 1500 C may not necessarily apply to higher temperature ranges. The oxidation resistance of this composite and the optimization of the composition at more than 2000 C are also under investigation.              2. Experimental procedure  2.1.  Samples preparation  Four kinds of ZrB2-SiC-BN composites with various powder compositions (vol%) were prepared (see Table 1). Commercially available ZrB2 powder (>99% purity, an average particle size of 3 \\u242em, Northwest Institute for non-ferrous metal research, China), SiC powder (>98.5% purity, an average particle size of 1.5 \\u242em, Weifang Kaihua Micro-powder Co., Ltd., China) and h-BN powder (>99% purity, 4 \\u242em, Chem Factory, Beijing, China) were used. Powders were mixed and ball-milled for 12 h in a polyethylene bottle charged with ethanol using ZrO2 balls. Solvent was then removed using a rotary evaporator. The dried powder mixtures were sintered by hot-pressing in an argon atmosphere at 1800 C and 23 MPa pressure for 1 h in graphite dies coated with pyrolytic BN.     2.2. Characterization  Bulk density and theoretical density were measured and assessed by the Archimedes method and the rule-of-mixture, respectively. Phase composition was identified by X-ray diffraction (XRD; Rigaku, Dmax-rb) using Cu K␣ radiation. The microstructure was characterized by field emission scanning electron microscopy (SEM; S4700, Hitachi, Tokyo, Japan) and  Table 1  Composition of the prepared composite samples.  Sample  ZrB2  ZS0  ZS1  ZS2  ZS3  80%  56%  43.1%  40%  SiC  20%  14%  26.9%  25%  BN  0%  30%  30%  35%  ZrB2 :SiC:BN  4:1:0  4:1:2.14  4:2.5:2.78  4:2.5:3.5  chemical compositions were evaluated by energy-dispersive X-ray spectroscopy (EDS; Phoenix, EDAX, Mahwah, NJ). Flexural strength was tested in a three-point configuration (3 mm × 4 mm × 36 mm chamfered bars), with a 30 mm span and a crosshead speed of 0.5 mm/min. Fracture toughness was evaluated by a single-edge notched beam test with a 16 mm span and a crosshead speed of 0.05 mm/min using 2 mm × 4 mm × 22 mm test bars. Hardness was determined by Vickers indentation (Model HVS-5, Laizhou Huayin Experimental Instrument Limited Company, China) using a diamond indenter with a load of 98 N for 15 s.  2.3. Oxidation tests     Specimens were cleaned in an ultrasonic bath in acetone before oxidation. The isothermal static oxidation tests were conducted in an electrical furnace at temperatures of 1100, 1300 and 1500 C in air with interruptions in the tests in order to measure weight (to an accuracy of 0.0001 g) change at fixed times. The specific weight change was calculated according to the mass change per surface area. The oxidation resistance of specimen was also tested by a Netzsch STA449C thermogravimetric analyzer. The mass changes were followed at a rate of 5 /min up to 1500 C with an 2 h isothermal hold in a flowing air (50 ml/min).        3. Results and discussion  3.1. Mechanical properties and machinability  Fig. 1 shows the XRD results of sintered ceramics containing various contents of BN. All samples contained the initial phases of ZrB2 , SiC and BN, except the sample ZS0 in which no BN existed. No new phase was formed during hot-pressing and sintering. Therefore, no chemical reactions occurred under the experimental condition, which benefits the formation of weak interfaces between the boundaries of ZrB2 , SiC and BN grains.  Fig. 1. XRD patterns of ZrB2 based composites.  \\x0c', 'H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  1037  Table 2  Density and mechanical properties of ZrB2 -SiC and ZrB2 -SiC-BN composites.  Sample  Composition (vol%)  Apparent density (g/cm3 )  ZS0  ZS1  ZS2  ZS3  80%ZrB2 + 20% SiC 56%ZrB2 + 14%SiC + 30%BN 43.1%ZrB2 + 26.9%SiC + 30%BN 40%ZrB2 + 25%SiC + 35%BN  5.129  4.102  3.775  3.735  Relative density  Flexural strength  (%)  93.0  90.3  90.6  92.6  (MPa)  281  301  317  378  Fracture toughness (MPa m1/2 )  4.8  3.5  3.7  4.1  Vickers hardness  (GPa)  15.9  5.9  5.6  52  Results of density and mechanical properties are listed in Table 2. An increase in the flexural strength of ZrB2-SiC composites doped BN was found compared to that of ZrB2-SiC composite without BN. This mainly results from the fact that the h-BN crystals were homogeneously dispersed around the matrix grains of ZrB2 and SiC during sintering (as shown in Fig. 2), which limits the grain growth and improves their flexural strengths. It is assumed that the soft h-BN particles with layeredstructures could relax stress and absorb energy at the crack tip through microcracking or crack-particle interactions, then prevent the main crack from extending which should be propitious to improve fracture toughness.13-15 However, compared to ZrB2-SiC, the fracture toughnesses of all ZrB2-SiC-BN composites decreased in the study. On the other hand, Table 2 shows that the hardness of the ZrB2-SiC-BN composite decreased greatly with 30 vol%BN additive compared to pure ZrB2-SiC. Hardness is an important indicator for ceramic machinability. Generally, a lower hardness leads to an improved machinability. Fig. 3 shows a hole made by cemented carbide drills on the ZS2 specimen. It can be seen that the ZrB2-SiC-BN composite is successfully machined. However, due to high hardness, the ZS0 specimen without BN additive cannot be machined using such drills. As stated above, the layered structure of BN resulting in a weak  interface at the ZrB2-SiC-BN grain boundaries is the main reason for the improvement of the machinability, which can enhance the crack deflection and avoid the catastrophic failure of the material during drilling. Fig. 4 shows the fracture surface of specimens for a test of fracture toughness. It can be seen that abnormal grain growth occurs in the ZS0 specimen with a main fracture model of transgranular fracture. For the ZS1, ZS2 and ZS3 specimens, fractures propagate parallel to the layer crystals because BN grains possess a layered crystal structure and are readily delaminated due to its low cleavage energy. Crack deflections, branching and blunting during machining of layered crystal BN are beneficial to prevent macroscopic fractures from propagation beyond the local cutting area, which lead to fracture modes dominated by the intergranular fracture. This phenomenon confirms the formation of weak ZrB2-SiC-BN interfaces by the addition of BN and is the main reason for the improved machinability of this composite.  3.2. Oxidation resistance  3.2.1. Thermal gravimetric analysis (TGA)  Fig. 5 shows the mass changes of the four specimens. It is shown that there is a similar tendency as the temperature below  Fig. 2. Cross-sectional SEM micrograph from polished section of ZS2 compos ite.  Fig. 3. Demonstration of the prepared machinable ZrB2 -SiC-BN ceramic composite using cemented carbide drill.  \\x0c', '1038  H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  Fig. 4. SEM micrographs of the fracture cross-sections of samples ZS0 (A), ZS1 (B), ZS2(C) and ZS3 (D).     1100 C and no significant weight gain was observed for the four specimens. While the weight gain rates of all the four samples increase abruptly as the temperature is up to 1500 C, which indicates an accelerated oxidation. Moreover, the weight gain rate of ZS0 is the fastest among the four samples. During the           isothermal oxidation stage at 1500 C, the weight gain rate slows down compared to those during heating up or anisothermal stage from 1100 to 1500 C, which results from the formation of oxide films on the surfaces. The formed oxide film actually acts as a barrier for further diffusion of oxygen into the fresh interface of ZrB2-SiC-BN. It is interesting that the addition of BN restrains the oxidation during not only heating up stage but also the high temperature isothermal oxidation stage (1500 C for 120 min), especially for ZS1 and ZS2 with medium contents of BN.     3.2.2. Oxidation resistance of the composites in static air environment     The isothermal oxidation resistances for all the specimens were studied at 1100, 1300 and 1500 C. The specific weight changes versus oxidation time are given in Fig. 6a-c. In general, the weight gain rates of all samples increase with the temperature and the specific weight change with time basically follows a parabolic oxidation law. The latter implies that the oxidation kinetics is controlled by transport of oxidant through the growing oxide film. At 1100 the highest  C, the sample ZS3 containing 35 vol%BN presents rate of specific weight gain. However, minimum     Fig. 5. TGA oxidation of ZrB2 based composites in air up to 1500     C.  \\x0c', 'H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  1039  Fig. 6. Weight change of composites at (A) 1100     C, (B) 1300     C and (C) 1500     C.  all processes of oxidation may proceed through the following steps:        C to 1100  the temperature from 600  (a) At C. The generated liquid B2O3 would heal the cracks in the oxide scale, leading to the partially or completely sealing of the cracks depending on the formation rate of B2O3 from the reaction (1) and (3) and the viscosity of B2O3 liquid.16           specific weight gain after 64 h, 0.0020 g/cm2 , was obtained for ZS2 (43.1%ZrB2 + 26.9%SiC + 30%BN), which is slightly lower than that of ZS0 (80%ZrB2 + 20%SiC), 0.0023 g/cm2 . At 1300 and 1500 C, the highest rate of specific weight gain is ZS0, and ZS2 presents the lowest oxidation rate. The weight gain for ZS0 and ZS2 after oxidation at 1500 C for 5 h were 0.0714 and 0.0268 g/cm2 , respectively. Therefore, it is concluded that the specimen ZS2 exhibits the best oxidation resistance from 1100 to 1500 C. For the modified composites, the expected main reactions describing the oxidation process are as follows: 2ZrB2 (s) + 5O2 (g) = 2ZrO2 (s) + 2B2O3 (l) 2SiC(s) + 3O2 (g) = 2SiO2 (s) + 2CO(g) 4BN(s) + 3O2 (g) = 2B2O3 (l) + 2N2 (g) B2O3 (l) = B2O3 (g) SiO2 (s) + xB2O3 (l) = B2O3 ·xSiO2 (l) B2O3 ·xSiO2 (l) = B2O3 (g) + xSiO2 (s)  (1)  (6)  (4)  (2)  (3)  (5)  The reaction (1)-(3) lead to weight gains, and the reactions (4) and (6) lead to weight loss. The weight change of sample is accumulative results of reactions (1)-(4) and (6). The over Fig. 7. XRD patterns of ZS0 (A) and ZS2 (B) oxidized at 1500     C for 1 h.  \\x0c', '1040  H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  Fig. 8. Cross-sectional micrographs of ZS2 after oxidation at 1500     C for 2 h.        (b) During the temperature from 1100  C to 1300  C. Large amount of volatile B2O3 would form. And very important, the formation of SiO2 derived from the oxidation of SiC becomes significant, reaction between SiO2 and B2O3 leads to a stable borosilicate glass. The borosilicates would act as a protective layer to reduce oxidation rate more effectively than B2O3 due to lower volatility and smaller oxygen diffusivity.17-19 The sample ZS3 gets the highest rate of specific weight gain, which is attributed to the largest content of BN (35 vol%). The formation of B2O3 from the oxidation of BN shows the obvious weight gain. For the sample ZS2,  it contains an appropriate amount of BN (30 vol%) compared with the sample ZS0, so more B2O3 was generated, and more easily spread in the material surface for oxide film. Moreover, the sample ZS2 contains a higher proportion of silicon carbide than ZS0 and ZS1, therefore more silicon oxide forms in ZS2 than that does in ZS0 and ZS1. The higher SiO2 content in borosilicate glass, the higher the viscosity and melting point of borosilicate glass are, which can more effectively cover and protect the surface of samples. Therefore, the oxidation resistance of ZS2 with a suitable amount of BN additive is better than other samples.  \\x0c', 'H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  1041  Fig. 9. Cross-sectional micrographs of ZS0 after oxidation at 1500     C for 2 h.        (c) With increasing temperature, 1300-1500  C. The viscosity of borosilicate glass decreases with increasing temperature, which benefits a healing of cracks and the diffusion velocity of oxygen. When temperature is up to 1500 C, the initially formed silica-enriched glass will be gradually lost in response to the reaction (6) due to the substantive volatilization of B2O3 . While SiO2 is significantly less volatile than B2O3 at these temperatures (the vapor pressure of B2O3 is 105 times higher than that of SiO2 at 1500 C),20 the remaining silicon oxide may react with zirconium oxide to generate a new anti-oxidation coating, zirconium silicate, which was confirmed by phase analysis in the study for the first time.        Fig. 7 shows the XRD patterns of the surface coatings of ZS2 and ZS0 oxidized at 1500 C for 1 h. Monoclinic ZrO2 and tetragonal ZrSiO4 were observed. Since the B2O3 and borosilicate are amorphous, some undetected B2O3 probably may remain dissolved in the SiO2 . The presence of ZrSiO4 , presumably derived from the chemical reaction between ZrO2 and SiO2 (reaction (7)), stabilizes SiO2 and inhibits the volatilization of silica. Besides these, ZrSiO4 has high melting point and vis cosity, it can efficaciously cover the sample surface and seal the cracks, which effectively limits the inward transport of oxygen and correspondingly enhances the resistance to oxidation. ZrO2 + SiO2 = ZrSiO4  (7)     Figs. 8 and 9 show SEM results for the oxidized ZS0 and ZS2 after oxidation at 1500 C for 2 h, respectively. It is noticeable that the oxide scales of both ZS0 and ZS2 consist of two distinct layers and the oxide layer of the specimen ZS0 (214 \\u242em) is thicker than that of ZS2 (169 \\u242em), which also indicates that the oxidation resistance of ZS2 is better than that of ZS0. The outer layers of both them are identified as ZrSiO4 , according to the combination of the EDS and XRD. In the inner layer, it is observed that white zirconia particles as a skeleton distribute in grayer zirconium silicate. The ZrO2 does not enhance the oxidation protection, but may provide mechanical integrity and strength like a framework for the liquid glass. At the same time, cracks were found in oxide layer during the quenching process, which attributes to the coefficients of thermal expansion mismatch between the oxide layer and matrix. EDS shows that zirconium, oxygen and silicon are present as the primary elements in the oxidized layer. Although quantitative  \\x0c', '1042  H. Wu, W. Zhang / Journal of the European Ceramic Society 30 (2010) 1035-1042  analysis for B is not possible by EDS since B is beyond the limit of the detection capability, a small amount of B element may still exist as borosilicate, though a previous SIMS investigation showed the amount being minimal.20 It is worth noting that the outer layer of ZS2 appears compact and adherent, while that of ZS0 is discontinuous and incompact. Furthermore, the thickness of ZS2 is remarkably lower than that of ZS0. This is due to the fact that the formation of this layer is mostly dependent on the SiC content. The silicon carbide content of ZS2 is higher than that of ZS0. A high SiC content is beneficial for the formation of ZrSiO4 , which can act as an effective barrier against the inward diffusion of oxygen.  4. Conclusions        (1) ZrB2 -SiC-BN ceramics were successfully prepared by hot-pressing under argon at 1800 C and 23 MPa pressure. With the addition of BN, the flexural strength of the ZrB2-SiC composite was improved, and the fracture toughness decreased slightly, but the hardness decreased enormously, and the machinability of this composite was improved noticeably. (2) Below 1300 C, the addition of BN significantly improved the oxidation resistance of ZrB2 -SiC ceramics due to the formation of ample borosilicates. At 1300 C and above, zirconium silicate deriving from the reaction between silica and zirconium oxide played as a major anti-oxidation coating, which could inhibit the diffusion of oxygen and protect the material underneath from further oxidation after evaporating of borosilicates. (3) The composite with components of 43.1%ZrB2 , 26.9%SiC and 30%BN showed excellent oxidation resistance up to 1500 C. A total weight gain as low as 0.0268 g/cm2 after oxidation at 1500 C for 5 h was observed. The addition of BN in the appropriate amount is implied as an effective method to simultaneously improve the flexural strength, machinability and the oxidation resistance of ZrB2-SiC ceramics.           Acknowledgements  Financial support from the Chinese Academy of Sciences under the Program for GF Basic Research (No. A1320070102) is gratefully appreciated.  References  1. Zimmermann, J. W., Hilmas, G. E., Fahrenholtz, W. G., Monteverde, F. and  Bellosi, A., Fabrication and properties of reactively hot pressed ZrB2 -SiC ceramics. J. Eur. Ceram. Soc., 2007, 27, 2729-2736.  2. Levine, S. R., Opila, E. J., Halbig, M. C., Kiser, J. D., Singh, M. and Salem,  J. A., Evaluation of ultra-high temperature ceramics for aeropropulsion use. J. Eur. Ceram. Soc., 2002, 22, 2757-2767.  3. Zhou, X.  J., Zhang, G.  J., Li, Y. G., Kan, Y. M. and Wang, P. L., Hot  pressed ZrB2 -SiC-C ultra high temperature ceramics with polycarbosilane as a precursor. Mater. Lett., 2007, 61, 960-963.  4. Monteverde, F., Bellosi, A. and Scatteia, L., Processing and properties of  ultra-high temperature ceramics for space applications. Mater. Sci. Eng. A, 2008, 485, 415-421.  5. Yang, F. Y., Zhang, X. H., Han,  J. C.  and Du, S. Y., Characteriza tion  of  hot-pressed  short  carbon fiber  temperature  ceramic  composites.  399.  reinforced ZrB2 -SiC ultra-high 2009, 395-  J. Alloys Compd.,  472,  6. Zhang, W. Y. and Gao, H., Microstructures, properties and fabrication of machinable ceramics. J. Synth. Cryst., 2005, 34, 170-173.  7. Wang, R. G., Pan, W., Jiang, M. N., Chen, J., Luo, Y. M. and Sun, R. F., Devel opment in machinable ceramics and machining technology of engineering ceramics. Bull. Chin. Ceram. Soc., 2001, 3, 27-35.  8. Xu, H. H. K. and Jahanmir, S S, Scratching and grinding of a machinable  glass-ceramic with weak interfaces and rising T-curve. J. Am. Ceram. Soc., 1995, 78, 497-500.  9. Zhang, G. J., Yang, J. F., Ando, M. and Ohji, T., Nonoxide-boron nitride  composites: in situ synthesis, microstructure and properties. J. Eur. Ceram. Soc., 2002, 22, 2251-2254.  10. Cho, M. W., Kim, D. W. and Cho, W. S., Analysis of micro-machining characteristics of Si3N4 -hBN composites. J. Eur. Ceram. Soc., 2007, 27, 1259-1265.  11. Li, Y. L., Zhang, J. X., Qiao, G. J. and Jin, Z. H., Fabrication and properties of machinable 3Y-ZrO2 /BN nanocomposites. Mater. Sci. Eng. A, 2005, 397, 35-40.  12. Kusunose, T., Sekino, T., Choa, Y. H. and Niihara, K., Fabrication and  microstructure of silicon nitride/boron Ceram. Soc., 2002, 85, 2678-2688.  nitride  nanocomposites.  J. Am.  13. Wang, R. G., Pan, W., Chen, J., Jiang, M. N. and Fang, M. H., Fabrication and  characterization of machinable Si3N4 /h-BN functionally graded materials. Mater. Res. Bull., 2002, 37, 1269-1277.  14. Li, Y. L., Qiao, G. J. and Jin, Z. H., Machinable Al2O3 /BN composite ceramics with strong mechanical properties. Mater. Res. Bull., 2002, 37,  1401-1409.  15. Wang, X. D., Qiao, G. J. and Jin, Z. H., Fabrication of machinable silicon  carbide-boron nitride ceramic nanocomposites. J. Am. Ceram. Soc., 2004, 87, 565-570.  16. Naslain, R., Design, preparation and properties of non-oxide CMCs  for  application in engines and nuclear Technol., 2004, 64, 155-170.  reactors: an overview. Compos. Sci.  17.  Jacobson, N. S. and Morscher, G. N., High-temperature oxidation of boron nitride: ii. Boron nitride layers in composites. J. Am. Ceram. Soc., 1999, 82,  1473-1482.  18. Baskaran, S. and Halloran, J. W., Fibrous monolithic ceramics: III. mechan ical properties and oxidation behavior of the silicon carbide/boron nitride system. J. Am. Ceram. Soc., 1994, 77, 1249-1255.  19. Guo, Q. G., Song, J. R., Liu, L. and Zhang, B. J., Relationship between oxi dation resistance and structure of B4C-SiC/C composites with self-healing properties. Carbon, 1999, 37, 33-40.  20. Rezaie, A., Fahrenholtz, W. G. and Hilmas, G. E., Oxidation of zirconium  diboride-silicon carbide at 1500 C at a low partial pressure of oxygen. J. Am. Ceram. Soc., 2006, 89, 3240-3245.     \\x0c']"
},{
  "_id": 77,
  "PDF": "Facile Precursor Synthesis of HfC-SiC Ultra-High-Temperature Ceramic Composite Powder for Potential Hypersonic Applications.pdf",
  "Text": "['Subscriber access provided by Kaohsiung Medical University  Article  A facile precursor synthesis of HfC-SiC ultra-high temperature ceramic composite powder for potential hypersonic applications  Niranjan Patra, and William E. Lee  ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.8b00781    Publication Date (Web): 08 Aug 2018  Downloaded from http://pubs.acs.org on August 15, 2018  Just Accepted  “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. 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However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.  • \\x0c', 'Page 1 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  A facile precursor synthesis of HfC-SiC ultra-high temperature  ceramic composite powder for potential hypersonic applications   Niranjan Patra∗, William Edward Lee   Centre for Advanced Structural Ceramics, Department of Materials,  Imperial College London, London SW7 2AZ, UK   Abstract   Nano-sized powder composites of hafnium carbide/silicon carbide (HfC-SiC) were successfully   prepared by crosslinking and carbothermic reduction of a precursor. A novel high yield liquid   precursor was synthesized by reacting acetylacetone (acac) modified hafnium tetrachloride (HfCl4)   with   tetraethyl orthosilicate (TEOS) and hydroquinone (HQ). The synthesized materials were   investigated by XRD, FT-IR, carbon analyzer, Raman, TG/DTA, SEM, TEM and EDS analysis.   The resulting precursor pyrolysed at 1500 oC for 3 h in argon converted into nanostructured ultra high temperature composite powder (containing both HfC and SiC) with a trace of free carbon. The   polymer-to-ceramic   transformation was   investigated by TG/DTA revealing   that carbothermic   reduction of HfC and SiC started at 1225 and 1450 oC respectively with a total yield of 64 % at   1500 oC. XRD revealed the powder composite was composed of HfC and β-SiC with a negligible   amount of oxide impurities. SEM and TEM images reveals the Hf, Si and C were homogeneously   distributed in the sample at submicron scale with an average grain size of less than 50 nm for both   HfC and SiC. The result obtained by this synthesis approach is believed to be a promising candidate   for hypersonic applications.   Keywords: Precursor; HfC-SiC; Ultra-high   temperature   ceramic;   composite; Carbothermal   reduction; Microstructure.   ∗Corresponding author: Department of Materials, Imperial College London,  London SW7 2AZ, UK.  Tel: +44 020 7594 1170, Fax: 020 75948175  E-mail address: patraji@gmail.com; drpatra.niranjan@gmail.com (N. Patra).   ACS Paragon Plus Environment  1                                                               \\x0c', 'ACS Applied Nano Materials  Page 2 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  1. Introduction   Ultra-high   temperature   ceramic   (metal   carbides)   have   numerous   potential   structural   applications like wing leading edges, nose tip of hypersonic re-entry vehicles, rocket nozzles,   thermal protection system which always involve harsh environments1,2. This is due to their excellent   combination of properties such as high melting point, excellent thermal and electrical conductivity,   good mechanical properties and chemical   inertness3-7. Among   the different   transition metal   carbides, hafnium carbide (HfC) is one of the most promising ultra-high temperature ceramic   (UHTC) having a melting point of ~3900 oC which has attracted much attention due   to its   outstanding physicochemical properties such as high hardness (26.5-32.9 GPa), high modulus of   elasticity (424 GPa), high electrical and thermal conductivity (20 W/mK) and chemical resistance in   harsh environments3,7,8,9. However, despite its suitable properties, sintering of monolithic HfC is   difficult due to its low self-diffusion coefficient and strong covalent bonding which means it   requires high temperature, usually above 2000 oC10. Such high temperature sintering can induce   rapid grain growth and coarse microstructures leading to poor mechanical properties11. HfC as well   as other transition metal carbides, borides and nitrides suffers poor stability in high temperature   oxidative atmosphere12,13. To overcome these problems, silicon carbide is generally introduced in   metal carbides (MC) (M=Hf, Zr, Ta) to produce MC-SiC composites which have been shown to be   more suitable in inhibiting the grain growth leading to improved mechanical properties compared to   single phase ceramic materials13-20,21. Additionally, the combination of passivating character of SiC   and excellent physico-chemical properties of HfC potentially generates a high-performance   composite able to endure aggressive environment21,22.          MC-SiC are conventionally prepared by solid state sintering of metal carbide and silicon   carbide powders19,22. However, sintering is difficult due to the coarse scale of the particles which   requires high temperature, pressure and longer production time19. The polymer-derived precursor is   a facile and promising approach for preparing dual phase ceramic nanocomposites with controlled   phase composition and microstructures and consequently with   improved properties23-25. This   ACS Paragon Plus Environment  2     \\x0c', 'Page 3 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  process reduces the kinetic barrier between the formed metal oxide and the carbon particles created   during precursor pyrolysis due to the homogeneous atomic scale dispersion of reactants. Increasing   the contact area of nano-grains allows carbothermic reduction between the metal oxide and carbon   particle to take place at lower temperature and in a shorter time26-29 compared to the conventional   solid state method which require temperature above 1800 oC4. Another advantage of polymer-  derived precursors over the conventional solid state route is the ability to form fibres and thin   films30,31.         The present study is focused on the synthesis of a precursor for pure, high yield binary solid   carbide mixture of HfC and SiC using combination of solution and carbothermic reduction enabling   fabrication of C/UHTCs composites by precursor   infiltration and pyrolysis. The polymer-to-  ceramic conversion and nano/microstructural evolution of the prepared HfC-SiC powder reveal a   novel approach to prepare ultra-high temperature composite powder from a high yield flexible   precursor.   2. Experimental          Hafnium tetrachloride (HfCl4 Mw=320.30 g/mol, 98% assay), acetylacetone (acac) (C5H8O2   Mw=100.12, 99.6% assay), hydroquinone (HQ)    (C6H6O2 Mw=110.11, ≥99% assay), tetraethyl   orthosilicate (TEOS) (SiC8H20O4 Mw=208.33, 98% assay) and anhydrous ethyl alcohol (C2H6O,   ≤0.005% water) obtained from Sigma-Aldrich Company Ltd, Gillingham, Dorset, UK were used   as-received without further purification for the synthesis of HfC-SiC precursors.         To synthesize HfC-SiC hybrid precursor, first hafnium tetrachloride (1M) was refluxed with   acetylacetone (2M) (Hf/acac) at 70 oC for 1 h using anhydrous ethanol as solvent to form a hafnium   metal diketonate conjugate structure as presented below.   ACS Paragon Plus Environment  3       \\x0c', 'ACS Applied Nano Materials  Page 4 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Acetylacetone functionalization helps to prevent precipitation of hafnium metal ion during rapid   hydrolysis and improve the solubility of metal ions in organic solvents. Subsequently hydroquinone   was added to the Hf/acac solution and refluxed for 5 h. Similarly, for the SiC, TEOS reacted with   HQ at 60 oC for 3 h using ethyl alcohol as solvent. The two precursor solutions were then mixed   together for another 1 h. A transparent orange solution was obtained at the end of the reaction and   the excess solvent was removed by a rota vapour (R-20 BUCHI Rotavapor, Switzerland) resulting a   gum like solid. The sample was further dried at 110 oC for 12 h to remove entrapped solvent   completely and crushed to powder in an agate mortar. Finally, pyrolysis was carried out at 1500 oC   for 3 h in controlled flow of argon (99.99% purity, 50 mL/min) in a tube furnace in order to avoid   oxidation. Two different HfC-SiC precursor   samples were prepared by varying   the   total   hydroquinone (HQ) concentration in order to study the optimum concentration to get pure HfC-SiC   composite powder. The first sample used a molar ratio of (Hf:acac:TEOS:HQ 1:2:1:1 M) herein   after referred to as HS1 and the other sample (Hf:acac:TEOS:HQ 1:2:1:2 M) herein after referred to   as HS2.          Fourier transform infrared spectra (FT-IR) of starting as-received materials and the synthesized   precursor powder after drying were acquired in the range of 4000-400 cm-1 in attenuated total   reflection (ATR) mode using a Nicolet iS10 instrument from Thermo scientific FT-IR spectrometer   Company Ltd, Waltham. USA. All spectra were baseline corrected and normalized thereafter to the   highest peak.          Thermal analysis was performed using a simultaneous thermogravimetry-differential thermal   analyzer (TG-DTA, NETZSCH STA 449 F1 Jupiter, Selb, Germany) in argon at a heating rate of   10 oC min-1 to study the polymer-to-ceramic conversion of the precursor. TG-DTA of pyrolyzed   product (HfC-SiC) was also carried out in air at a heating rate of 10 oC min-1 which provides   information about the oxidative stability, particle size, uniformity and free carbon content in the   samples26-29.    ACS Paragon Plus Environment  4     \\x0c', 'Page 5 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60        X-ray   diffraction   (XRD) was   performed   using   a   computer   controlled   diffractometer   (PANAlytical Xpert3 diffractometer, Cambridge, UK) using Cu Kα1 radiation (λ=1.540598 Å) to   identify the phases. The X-ray tube was operated at 40 kV, 40 mA with scanning step width of   0.030 (2θ). The crystallite sizes were determined from the broadening of the XRD diffraction using   the Debye-Scherrer equation. International centre for diffraction database (JCPDS Card No. 65-  0975, and 74-2307) were used in Xpert High Score Plus software to determine crystalline phases of   HfC and SiC respectively.         Carbon content was measured   in a carbon analyzer (Horiba EMIA-920V2 C/S analyzer,   Northampton, UK) using tungsten carbide (WC) (BCS-CRM No. 352/1) as the standard reference   of Bureau of Analysed Samples Ltd (BAS), UK32. Raman spectra were acquired in a Raman   spectrometer (Renishaw, Gloucestershire, UK) with a resolution of 4 cm-1 using  an incident   radiation from an Ar ion laser at λ=514.5 nm, coupled with an optical microscope with a 50X   objective lens. Acquisition times were on the order of 50 s. For all experiments three measurements   were carried out and averaged for consistency and reproducibility.        Morphology was examined in an SEM (JEOL JSM 6010 LA, Tokyo, Japan) using secondary   electron imaging (SEI) mode using an accelerating voltage of 20 kV coupled with an energy   dispersive X-ray (EDS) analyzer fitted with a silicon drift detector (JOEL SDD) with an ultra-thin   window. SEM Specimens were prepared using aluminium stubs of size 12.5mm dia, 10mm high as   sample holder. A double sided sticky carbon tape made of spectroscopically pure carbon is used to   glue the sample to the stub which also avoids undesirable background radiation from aluminium   stubs. A pinch of sample powder was taken using a spatula and put it on the sticky carbon tape in   the aluminium stubs and pressed using a glass slides to make it a smooth flat surface. The sample   was then coated with 10 nm of gold for better conduction. Morphology of the fine particles was also   characterized by transmission electron microscopy (TEM, JEOL JEM 2100F, Tokyo, Japan) with   an acceleration voltage up to 200 kV using bright-field imaging. EDS mapping was performed   using INCA EDS 80 mm X-Max detector system (Oxford instrument, Abingdon, UK) coupled with   ACS Paragon Plus Environment  5     \\x0c', 'ACS Applied Nano Materials  Page 6 of 31  the TEM. For TEM specimen, a pinch of powder was sonicated using ethanol as solvent. A small   drop of (2 µL) solution was dropped on to a carbon coated TEM copper grid followed by drying.   3. Results and discussion         Figure 1a, b and c shows   the normalised FT-IR absorption spectral band of synthesized   precursors (HS1 and HS2) dried at 110 oC, hafnium-(acac)diketonate conjugate structure and   starting reactant materials (HQ, HfCl4, acac and TEOS) used for synthesis of the hybrid precursor to   determine the structural interactions and changes of organic polymer ligands with the inorganic   metal ions (e.g. Hf and Si). For HQ (Fig. 1c) shows more characteristics absorption bands in the   fingerprint region. The major bands observed at 3160, 3030 and 1517/1464 cm-1 were assigned to   the O-H, C-H and C C stretching; bands at 1204/1096 and 824/757 cm-1 correspond to C-O   stretching and C-H out-of-plane bending on the aromatic ring respectively33. For HfCl4 (Fig. 1c),   the band at 3328 cm-1 is assigned to O-H stretching of physisorbed molecular water while a low   intensity band at 3153 cm-1 is assigned to C-H stretching. The stretching band at 1590 cm-1 is   assigned   to   the bending   vibration of water molecules   coordinated   to   the   cations34. For   acetylacetone, the major band observed corresponds to the characteristic groups of C O, C C at   1730 and 1707 cm-1. For TEOS, the band at 1073 cm-1 corresponds to Si-O-C; that at 465 cm-1   could be assigned to the O-Si-O bond; the strong band at 783 cm-1 corresponds to the Si-O band;   and finally the weak band at 1290 cm-1 corresponds to Si-CH3 35. Fig 1b shows the FTIR spectrum   of the conjugate structure formed by reaction of hafnium alkoxide with acetylacetone. The C=C   (enol form), C=O (keto form) and C-CH3 absorption bands shifted to lower wavenumber at 1574,   1538 and 1275 cm-1 which confirms hafnium alkoxide (HfCl4) reacted with acetylacetone (acac) to   form hafnium diketonate conjugate structure (as suggested in the inset structure in fig 1b). The   diketonate conjugation contributed to the stability of the liquid precursor by preventing metal ion   precipitation during rapid hydrolysis conferring a long shelf life on the prepared precursor. For the   as-synthesised hybrid precursor (HS1) (Fig. 1a), the intensity of stretching vibrations band at 1595    6   ACS Paragon Plus Environment  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60      \\x0c', 'Page 7 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  and 1720 cm-1 for C=C and C=O from acetylacetone diminishes indicating the precursor could be   dehydrogenized during heating27. The O-H stretching band appears at 3290 cm-1. FT-IR also   revealed the absorption bands that can be assigned to the presence of Si-CH3 bonds (1282 cm-1)   from HS1 and HS2. Absorption bands at 433 cm-1 and 512 cm-1 confirm the presence of Hf-O in   the main chain of precursor resulted from hydrolysis and condensation reactions. In addition, the   peaks at 1282 and 1095 cm-1 are attributed to Si-CH3. The peaks at 826 cm-1, 1040 cm-1, 1094 cm-  1 are attributed to Si-C, Si-O-Si and Si-CH3 in the precursor, respectively.    Fig 1. FT-IR spectra of (a) as-synthesized sample (HS1 and HS2) dried at 110 oC (b) hafnium-  (acac)diketonate conjugate (c) starting reactant materials.   ACS Paragon Plus Environment  7       \\x0c', 'ACS Applied Nano Materials  Page 8 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60        Fig 2 shows TG/DTG/DTA of the precursor from RT to 1600 oC in argon. The pyrolytic   transformation of hybrid organic-inorganic precursor   to an   inorganic network also known as   ceramization or mineralization involves two major steps: carbonization and carbothermal reduction.   TG/DTG/DTA reveals that the precursor shows five regions of mass loss. The first three steps show   evaporation of residual solvent, moisture, unreacted organic chains and carbonization while the last   two corresponding to the carbothermal reduction leading to an overall mass loss of 43 wt% at 1600   °C: Step (1) RT to ≈200 °C (≈8 wt%) due to the evaporation of residual solvent and water from   polycondensation; Step (2) 200 to ≈673 °C (≈17 wt%) due to the decomposition of organic chains   to form carbon and dehydration of some hafnium alkoxide; Step (3) 673 to ≈1233 °C (≈3 wt%) due   to evaporation of CO2 during HfO2/SiO2 formation, Step (4) 1233   to ≈1461 °C (≈6 wt%),   corresponding to carbothermal reduction to form HfC by loss of CO and CO2; Step (5) 1461 to   ≈1585 °C (≈8 wt%) corresponding to the carbothermal reduction to form SiC27,36. The total HfC-  SiC yield at 1600 oC is 65% excluding the moisture and entrapped solvent which is around 8% at   200 oC. From the TG/DTG results of HS in argon, the overall ceramization process can be described   as below.                         HS(s)                                  HfO2(s)  +  SiO2(s)  +  C(s)  + Volatiles(g)              (1)                                        HfO2(s) + 3C(s)                                   HfC(s)  + 2CO(g)                               (2)                                        SiO2(s) + 3C(s)                                               SiC(s) + 2CO(g)                           (3)        HfCl4, TEOS and HQ in precursor solution are used as hafnium, silicon, and carbon source,   respectively. The HS is converted to HfO2 and SiO2 at 800 oC and converted into HfC and SiC   started at 1225 and 1450 oC respectively. The precursor transformed completely to crystalline HfC   and SiC at 1500 oC. The onset of low carbothermic reaction temperature could be responsible due to   the intimate mixing of reactants26-28.   ACS Paragon Plus Environment  8       \\x0c', 'Page 9 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 2. TG, DTG and DTA of as-synthesized hybrid precursor (HS1) in argon.         The phase composition of the synthesized powders (HS1 and HS2) containing HfC and SiC heat   treated at 1500 oC for 3 h at constant argon flow rate of (50 ml/min) was investigated via XRD   analysis shown in Fig. 3a, b. Peaks of well crystallized HfC and β-SiC are clearly identified for   both HS1 and HS2 precursors. Fig 3b shows the magnified peaks of SiC for clarity. Peaks at 33.5,   38.8, 56, 66.8, and 70.2 were identified corresponding to HfC phase, while the other four peaks at   2θ of 35.5, 41.4, 60 and 70.3 were found consistent with β-SiC. No other peaks were observed in   the pattern. The crystallite sizes of HfC and β-SiC calculated from the Debye-Scherrer equation   (Dhkl=kλ/βcosθ) were 47 nm and 27 nm, respectively. The reduced nano crystallite size could be   due to the inhibition of crystallization between HfC and SiC crystalline phases and homogeneous   element distribution from the intrinsic properties of the precursor20.   ACS Paragon Plus Environment  9         \\x0c', 'ACS Applied Nano Materials  Page 10 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 3(a). XRD of composite powders HS1 and HS2 heat treated 3 h at 1500 oC in argon, (b)  magnified pattern of curves (a) for the sake of clarity for β-SiC peaks (gray colour highlights).   Table 1 shows the results of carbon analysis of the HfC-SiC composite powder (HS1 and HS2)   samples revealing the HS1 sample contains the least total carbon content (8.35 wt%) which is   slightly higher (~1.8 wt%) compared to the stochiometric sample of HfC which is 6.55 wt%   (99.9%, 100nm)37. The HS2 samples show an undesirably high total carbon content of 20.4 wt%   which would detrimentally affect sintering by releasing CO2 during high temperature (1800- 2300   oC)  sintering thus making the material porous38. The high total free carbon content in HS2 is due to   the excess of hydroquinone (HQ). The slightly greater free carbon content (~1.8 wt%)  in HS1   could prove to be beneficial during sintering as it will help prevent the grain growth at high   temperature in agreement with the earlier study39.   ACS Paragon Plus Environment  10             \\x0c', 'Page 11 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Table 1: EDS and carbon analysis of HfC-SiC powders calculated using WC (BCS-CRM No.   352/1)32.   Sample   HS1   HS2   Hf   64   57   Si   27   23   Elements (wt%)   C   8.6   19.8   O   0.3   0.2   Total C(carbon  analyzer)   8.4±0.4   20.4±0.4         The evolution of free carbon in the ceramic samples was identified by Raman spectroscopy.   Figure 4 shows the Raman spectra of HS1 and HS2 heat treated at 1500 oC for 3 h. Two prominent   peaks appeared at 1337 and 1586 cm-1 corresponding to D (disordered-induced) and G (graphite)   peaks observed in the case of free carbon40. The peak at 1337 cm-1 is attributed to disordered carbon   with lattice defects40 while the other at 1586 cm-1 attributed to in-plane stretching (E2g) symmetric   vibration in graphite41. Raman spectra show weak D and G bands. The degree of disorder of (ID/IG)   was rationalized on the basis of interdefect distance (LD) which could be defined as ID/IG=C(λ)/LD 42.   With low carbon, the integrated intensities of ID/IG decrease significantly indicating LD becomes   larger. Week D and G band for low HS1 indicates almost pure HfC and SiC with a quite low free   carbon content.   ACS Paragon Plus Environment  11                     \\x0c', 'ACS Applied Nano Materials  Page 12 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 4. Raman spectra of the samples (HS1 and HS2) after heat treatment at 1500 oC in argon.         Thermogravimetric analysis of carbide oxidation in air flow has been used to investigate the   particle size, distribution and homogeneity of the synthesized powder revealing that transition metal   carbides show a lower onset and small range of oxidation temperature when the particles sizes are   nanometer scale and homogeneous28,29. Fig. 5 shows the TG and DTA curves of the HS1 and HS2   samples in air. TG curves show two regions of oxidation mass gain as well as two regions of   decomposition mass loss in the composite powder. For both HS1 and HS2, the first region of mass   gain corresponds to oxidation of HfC whereas the second region of mass gain corresponds to   oxidation of SiC. The onset of oxidation for HfC started at ~220 oC with a maximum increase in   mass of 13.7 and 10.4 wt% at 624 and 582 oC for HS1 and HS2 respectively. The increase in mass   is due to the oxidation of HfC to form hafnium oxide and carbon. For commercial HfC powder of   average particle size (<1.5 µm), the onset of oxidation observed by Epshteyn et al.43 was 300 oC   with a mass gain of 11.1 wt%. The first step of oxidation mass gain of HS1 and HS2 samples starts   at   low   (230   oC)   and narrow   temperature   range   (230-624   oC)   indicating   submicron   and   homogeneous grain size distribution. The low onset of oxidation (230 oC) for HfC in both HS1 and   ACS Paragon Plus Environment  12         \\x0c', 'Page 13 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  HS2 is due to the high reactivity of the nanosized HfC particles27. The oxidation mass gain for HfC   component is higher for HS1 (13.7 wt%) than HS2 (10.4 wt%). HS1 reveals a large amount of HfC   (64 wt%) with a trace of free carbon whereas HS2 contains 20 wt% of free carbon. The mass loss of   5.8 and 19.3 wt% for HS1 and HS2 started after the oxidation which is due to the combustion of   free carbon to form CO2 12. The total carbon content of 5.8 and 19.3 wt% also nearly agrees well   with the results of carbon analysis. The second onset of oxidation temperature for SiC in HS1   observed at 750 oC is lower than HS2 (900 oC) with total mass increase of 6.4 and 4.5 wt%. By   comparing the HS1 and HS2, it is obvious that the onset of oxidation for SiC component was   postponed with   increasing free carbon content which acted as a diffusion barrier   layer and   suppresses or postpone oxidation. The DTA curve revealed that there are three exothermic reactions   occurring during oxidation. These exothermic reactions are represented by peaks located at 565,   655 and 1189 oC for HS1 and 489, 845 and 1210 oC for HS2 samples respectively. The exothermic   peaks can be explained as follows; the lower temperature exothermic peaks (565 and 489 oC for   HS1 and HS2) are due to the oxidation of HfC to form hafnium oxide leading to increasing mass.   The second exothermic peaks (655 and 845 oC for HS1 and HS2) are associated with CO2 evolution   from carbon combustion. The third exothermic peaks (1189, 1210 oC for HS1 and HS2) are   associated with oxidation of SiC to form silicon dioxide. From the above results, the oxidation   process can be described as below:                         HfC(s) + 2O2(g)                                           HfO2(s) + CO2(s)             (4)                          SiC(s) + 2O2(g)                                      SiO2(s) + CO2(g)                   (5)   ACS Paragon Plus Environment  13       \\x0c', 'ACS Applied Nano Materials  Page 14 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 5. TG (solid line) and DTA (dotted line) of oxidation of HfC-SiC composite powder (HS1 and   HS2) in air.         An SEM image of HfC-SiC powder (HS1) pyrolyzed at 1500 oC is shown in Fig 6a. HfC-SiC   particles showed near spherical morphology with narrow size distribution with particles sizes in the   range of 50 to 100 nm. The slightly larger particles size which is greater than calculated crystallite   size from the Scherrer equation possibly arise from aggregation of smaller crystallites of HfC and   SiC. The crystallite sizes of HfC and β-SiC calculated Debye-Scherrer equation (Dhkl=kλ/βcosθ)   were 47 nm and 27 nm, respectively. The EDS shown in fig 6b confirms the presence of hafnium   and silicon in the synthesized powder with a trace of free carbon.   ACS Paragon Plus Environment  14         \\x0c', 'Page 15 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 6(a, b) SEM image and EDS of HfC-SiC powder composite prepared from HS1 precursor at  1500 oC         HfC-SiC particles (HS1) synthesized at 1500 oC were studied by TEM with respect to phase   composition and micro/nanostructure (Fig 7a). Bright-field images showed dark contrast HfC   particles and light contrast β-SiC particles. Fig 7a revealed that the nanostructured samples were   homogeneously distributed with near   spherical morphology   although with   some   faceting.   Additionally, graphite like free carbon was also detected in EDS, which is consistent with the   results of Raman, TGA in air and carbon analysis. The nanocrystals ranged in size from 25 to 50 nm   consistent with the average size estimated by the Debye-Scherrer formula. On the basis of TEM   results combined with SEM and XRD, the small particle size of HfC and SiC in precursor derived   composites with a trace of free carbon is largely due to crystallization of HfC and SiC. These results   indicate phase evolution of the HfC-SiC composites prepared by pyrolysis of the precursor was   highly flexible. Micro/nanostructure and low pyrolysis temperature (1500 oC) of these ceramic   composites suggest that this is a promising route for their fabrication for high temperature aerospace   structural applications in harsh environments.   ACS Paragon Plus Environment  15           \\x0c', 'ACS Applied Nano Materials  Page 16 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig. 7(a). Bright-field TEM image and (b) size analysis (average particle size) of HS1 precursor  pyrolyzed at 1500 oC         Precursor-derived-ceramic (PDC) solutions are well known for homogeneous atomic level   distribution of different reactant component26,28. EDS mapping of HfC-SiC particles from HS1   precursor heat treated at 1500 oC shown in Fig 8 reveals Hf, Si and C were well distributed at   submicron scale without any phase separation. Hence, the element homogeneity of the precursor-  derived-ceramics is clear.    ACS Paragon Plus Environment  16         \\x0c', 'Page 17 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig. 8 TEM image and element distribution maps of composite particles (HS1) obtained at 1500 oC  in argon (a) TEM image (mapping zone), (b) EDS layered image of mapping zone, (c) EDS graph  of particles and (d) EDS maps of Hf, Si and C.   Conclusions         We have developed a liquid precursor route to prepare hafnium carbide-silicon carbide (HfC-  SiC) ultra-high temperature ceramic composite powder. The precursors were prepared by reacting   acetylacetone modified hafnium tetrachloride with tetraethyl orthosilicate and hydroquinone. The   synthesized precursor has excellent solubility in organic solvents and is air stable. Pyrolysis of the   precursor at 1500 oC produces well crystallized particles of HfC and SiC of average particle size   less than 50 nm with a trace of free carbon (1.8%) revealed by carbon analysis, TGA and Raman   spectroscopy. The trace of carbon in the Hf-SiC composite powder will help improve oxidation   resistance as well as prevent grain growth during high temperature sintering. This polymeric liquid   precursor synthesis approach has excellent capability for precursor-infiltration and pyrolysis in   multi-component ceramic composites fabrication.   ACS Paragon Plus Environment  17           \\x0c', 'ACS Applied Nano Materials  Page 18 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Acknowledgement   The financial support provided by Office of Naval Research Global (ONRG), USA under contract   number N62909-13-1-N055 is gratefully acknowledged.   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ACS Paragon Plus Environment  22   1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60                        \\x0c', 'Page 23 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  A facile precursor synthesis of HfC-SiC ultra-high temperature   ceramic composite powder for potential hypersonic applications   A novel organic-inorganic hybrid precursor   synthesized   for ultra-high   temperature ceramic   composites of HfC-SiC having excellent ceramic yield and infiltration capability reported.    ACS Paragon Plus Environment  23           \\x0c', 'ACS Applied Nano Materials  Page 24 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 1. FT-IR spectra of (a) as-synthesized sample (HS1 and HS2) dried at 110 oC (b) hafnium-  (acac)diketonate conjugate (c) starting reactant materials.    207x213mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'Page 25 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 2. TG, DTG and DTA of as-synthesized hybrid precursor (HS1) in argon.    207x145mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'ACS Applied Nano Materials  Page 26 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 3(a). XRD of composite powders HS1 and HS2 heat treated 3 h at 1500 oC in argon, (b) magnified   pattern of curves (a) for the sake of clarity for β-SiC peaks (gray colour highlights).    146x103mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'Page 27 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 4. Raman spectra of the samples (HS1 and HS2) after heat treatment at 1500 oC in argon.    138x92mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'ACS Applied Nano Materials  Page 28 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 5. TG (solid line) and DTA (dotted line) of oxidation of HfC-SiC composite powder (HS1 and HS2) in air.    146x103mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'Page 29 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig 6(a, b) SEM image and EDS of HfC-SiC powder composite prepared from HS1 precursor at 1500 oC    117x54mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'ACS Applied Nano Materials  Page 30 of 31  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig. 7(a). Bright-field TEM image and (b) size analysis (average particle size) of HS1 presursor pyrolyzed at   1500 oC    190x142mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c', 'Page 31 of 31  ACS Applied Nano Materials  1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60  Fig. 8 TEM image and element distribution maps of composite particles (HS1) obtained at 1500 oC in argon   (a) TEM image (mapping zone), (b) EDS layered image of mapping zone, (c) EDS graph of particles and (d)   EDS maps of Hf, Si and C.    190x142mm (300 x 300 DPI)    ACS Paragon Plus Environment               \\x0c']"
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  "_id": 78,
  "PDF": "Formation of Oxide Scales on Zirconium Diboride-Silicon Carbide Composites During Oxidation- Relation of Subscale Recession to Liquid Oxide Flow.pdf",
  "Text": "['Formation of Oxide Scales on Zirconium Diboride-Silicon Carbide  Composites During Oxidation: Relation of Subscale Recession to Liquid  Oxide Flow  Sigrun N. Karlsdottir  w,z  and John W. Halloran  Department of Materials Science and Engineering, University of Michigan, Ann Arbor, Michigan 48104  The formation of oxide scales on zirconium diboride (ZrB2)- silicon carbide (SiC) composites oxidized at high temperatures (415001C) ZrB2 composites is proposed to form due to flow of boron oxide (boria) (B2O3)-rich borosilicate liquid through convection cells that form upon oxidation of the composite at high temper is  studied. Subscale  recession  found  in  oxidized  atures. The flow of the B2O3-rich liquid to the surface, with the subsequent loss of B2O3 to evaporation, explains the formation of a glassy silica-rich layer on the surface commonly reported in  the  literature. Also the outward flow of  the  liquid creates a  localized inward path for oxygen due  to lower  viscosity that  allows faster oxidation under the convection cells which creates  the subscale recession. Optical and electron micrographs of a ZrB2-15%SiC composite oxidized at 15501C are presented as evidence of flowing liquids. Micrographs of oxide scale defor mations are also presented, which are proposed to be related to  the formation of oxide scale features called convection cells. The subscale recession and oxide scale deformations of ZrB2- 15%SiC composites oxidized for 3 and 4 h at 15501C were studied with microstructure and chemical composition analysis.  I.  Introduction  THE  zirconium diboride  (ZrB2)-silicon HfB2-SiC composites oxidize to form a complex multilayer oxide scale at temperatures between 14001 and 17001C.1-3  carbide  (SiC)  and  Often the oxide scale features a silica (SiO2)-rich outer which lies over a subscale of crystalline zirconia (ZrO2), often with a columnar microstructure with SiO2 between the ZrO2 grains. Depletion of SiC from the virgin material under the ZrO2 scale has also been observed and reported.4,5 While frequently  layer,  observed, the mechanisms that form this complex scale are not understood in detail.6,7 The interpretation of this complex oxide  scale presents several puzzles. ZrO2 appears often as a columnar subscale and as a noncolumnar phase in the SiO2 layer. Depletion of SiC underneath the ZrO2 scale suggests formation of SiO2 liquid or SiO vapor under the ZrO2, but most of the SiO2 is found over the ZrO2. The SiO2 liquid can dissolve with boron oxide (boria) (B2O3) liquid (formed upon oxidation of ZrB2) forming a borosilicate surface layer (B2O3-SiO2). B2O3 is, however, largely absent, due to evaporation at higher temperatures. B2O3 has a vapor pressure of 233 Pa at 15001C8 while SiO2 has a vapor pressure of 3 \\x02 10\\x004 Pa.5 How can these complex oxide scale features be interpreted? al.9 proposed that  Recently, Karlsdottir  et  liquid flow of  B2O3-SiO2-ZrO2 formation of these scales, based on distinctive microstructural  (BSZ)  liquids plays an important role in the  features observed on the  external oxide  surface and in cross  section.  These features are called convection cells. Figure 1 shows an  example of these convection cells. The image shows a backscat tering electron (BSE) image of a surface of a ZrB2-15 vol% SiC specimen oxidized at 16001C for 30 min. The surface of the  oxidized specimen is covered with convection cells  forming a  pattern. The  convection cells have ZrO2 located in larger SiO2-rich ‘‘lagoons’’(gray area) with B2O3-rich patterns (dark contrast) surrounding the islands. The area  islands  (white area)  around the convection cells consists of a SiO2-rich glass (gray region) with small micrometer-sized ZrO2 dispersoids (white dots). The B2O3 flower-like patterns are visible in BSE images, but in stronger contrast when imaged by cathodoluminescence (CL). The ZrO2 islands have been proposed previously9 to have formed by precipitation during the evaporation of B2O3 from a BSZ liquid that rises through an outer SiO2-rich borosilicate layer and flows laterally by viscous fingering forming the B2O3-rich regions around the ZrO2 islands.9,10 The driving force is proposed to be the large volume increase upon oxidation of the bulk material due to the formation of solid ZrO2(s) and BSZ liquid.9 Figure 2 is a schematic of these convection cells and their formation.  These convection cell features had not been discussed by oth ers before, but close examination of cross-sectional micrographs  in the literature suggest  that  the features might have been ob served, but have not be interpreted. In this paper the interpre tation of oxide scale features of diboride-SiC composites will be  discussed. Optical and electron micrographs will be presented as  evidence of flowing liquids. Oxide scale deformation related to  the formation of convection cell will also be presented and dis cussed. Finally, subscale recession found in oxide scales formed  during oxidation of boride-SiC-based material are discussed in  connection to convection cell features.  II.  Experimental Procedure: Materials and Methods  ZrB2-15 Institute of Science  vol% SiC composite materials were  fabricated  at  and Technology  for Ceramics, National  Research Council (CNR-ISTEC) in Faenza, Italy, using methods presented elsewhere.11 Before testing, ca. 200 mm was  re moved from surface by diamond grinding (Omni Brade, TBW  Industries, Furlong, PA). This was to remove any heat-affected  zone  that could have formed during wire electrical discharge  machining (w-EDM)  (Ann Arbor Machine Model 1S15, Ann  Arbor, MI), which was used to cut the bulk material  into thin  sheets. The thin sheets of the ZrB2-15 vol% SiC material were then cut with a diamond saw (IsoMet 1000 diamond precision  s  saw, Buehler, Lake Bluff, IL) into small rectangular coupons with total surface area on average of ca. 1 cm2. They were ox idized at high temperatures in ambient air at temperatures between 15501 and 16001C for different times ranging from 1  2  to  4 h. The tests were performed either in a high-temperature box  furnace (SentroTech Corporation, Berea, OH) or in a tube-fur nace (Lindberg, Watertown, WI). The heating rate used was 131C/min with free cooling or with 131C/min cooling rate. In the  H.-J. Kleebe—contributing editor  w  Author to whom correspondence should be addressed. e-mail: nanna@umich.edu zPresent address: Department of Materials, Biotechnology and Energy, Center Iceland, IS-112 Reykjavik, Iceland.  Innovation  Manuscript No. 24490. Received March 31, 2008; approved July 10, 2008.  Journal  J. Am. Ceram. Soc., 91 [11] 3652 - 3658 (2008)  DOI: 10.1111/j.1551-2916.2008.02639.x  r 2008 The American Ceramic Society  3652  \\x0c', 'November 2008  Formation of Oxide Scales on ZrB2-SiC Composites  3653  Fig. 1.  Backscattered electron images of  the surface of a ZrB2-15 vol% SiC composite oxidized at 16001C for 30 min. convection cells that have ZrO2 islands (white) located in larger SiO2, lagoons (gray) with B2O3-rich patterns (black) surrounding the islands. (b) Higher magnification of a convection cell.  (a) Surface covered with  furnace  the  specimens were  supported by the  same material  the glass is suggestive of liquids flowing from the neighborhood  (ZrB2-SiC) crucible.  that was placed on an Al2O3 support  in an Al2O3  Chemical composition and microstructural analysis were done  on the surfaces and cross sections of the oxidized specimens using  bright field optical microscopy of  the as-oxidized surface, scan ning electron microscopy, BSE microscopy, and electron micro probe analyzer (EMPA). A Cameca SX100 was used for EMPA,  of the ZrO2 island. Figure 3(b) is a BSE image of the same field of view as Fig. 3(a). In BSE imaging, the ZrO2 islands appear in bright contrast. The SiO2-rich glassy regions appear in light gray contrast, and the B2O3-rich regions in darker contrast. The small dispersed ZrO2 particles on the SiO2 are highly visible. The arrangement of these small ZrO2 dispersoids is hard to understand without presuming that they were arranged by flowing liquids.  using well-characterized mineral standards for quantitative anal We present these images as evidence in support of the hypothesis  ysis of boron (B), oxygen (O), zirconium (Zr), and silicon (Si),  of convective flow, as illustrated in the schematic of Fig. 2.  and for imaging in the BSE and CL modes. The EMPA stan Now let us direct our attention to cross section of the oxide  dards and technique that were used are described in more detail elsewhere.9 The  the oxidized specimens were  sections of  cross  prepared for microstructural analysis by nonaqueous polishing procedures down to 1 mm finish. Specimens were coated with  carbon before microstructural and elemental analysis.  III.  Interpretation of Oxide Scale  Examination of  the  surface of  the oxidized samples provides  clear evidence of liquid flow. Figure 3(a) is an optical micrograph of the surface after oxidation for 4 h at 15501C. Optical  metallograph image in reflected light shows what appears to be  ‘‘islands’’ of ZrO2 in a film of borosilicate glass. These islands are assemblies of ZrO2 grains emerging from the once-liquid glassy surface. The darker regions of the glass are rich in SiO2, and are dark because of the relative transparency scatters little  of the incident light. Very small dispersed ZrO2 particles barely visible are on the surface of the SiO2. The cloudy features are subsurface B2O3-rich borosilicate. These borosilicate regions are turbid because of liquid-liquid phase separation in the glass  during cooling. The appearance of the turbid and clear regions of  Fig. 2.  Schematic of the convection cell  features seen on surfaces and  cross sections of ZrB2-SiC composites oxidized at temperatures between 15001 and 16001C. The schematic shows convection cells spread on the  surface with their ZrO2 islands (white) located in larger SiO2, (light gray) with B2O3-rich patterns (dark gray) surrounding the islands.  lagoons  scales. Figures 4(a)-(c) show BSE images of a cross section of a ZrB2-15 vol% SiC specimen oxidized at 15501C for 3 h in a tube-furnace in ambient air, with a heating rate of 131C/min and  free cooling. The images show significant deformations of  the  oxide scale. The SiO2 outer scale appears in the BSE image in darker contrast, while the ZrO2 subscale is in brighter contrast. The contrast between ZrO2 and ZrB2 is slight in this image. The higher magnifications, Figs. 4(b) and (c), show that the ZrO2 scale is deformed; it appears to be lifted up, like a blister. Inside  the ‘‘blister’’  is a glassy phase shown by EMPA maps and anal ysis to be rich in B2O3, SiO2, and with some ZrO2. Figures 5(a)-(e) show a BSE image of one of the deforma tions shown in Fig. 4(a) and the corresponding EMPA maps of  this area, showing the distribution of B, Si, O, and Zr. The dis tinction between the ZrO2 in the primary scale and the ZrB2 substrate can be made by comparing the zirconium image  Fig. 5(c) with the boron image Fig. 5(e) and the oxygen image  Fig. 5(b). Clearly the ZrB2 substrate and the ZrO2 primary scale are being separated by a liquid rich in O, Si, and B, with sig nificant distortion of the ZrO2 primary scale. Figure 6 shows the corresponding line analysis from these elemental maps, indicat ing that the glass inside the ‘‘blister’’  is rich in B2O3, and SiO2, calculated ternary phase  on  a  has  and  some ZrO2. Based diagram of a ZrO2-SiO2-B2O3 system, an isothermal section at 15001C, published previously by Karlsdottir et al.,10  it  is  presumed that this material  is the glass formed by cooling of a  BSZ liquid in equilibrium with ZrO2(s) (15501C).  at  this  temperature  The driving force for  these deformations  is  likely the very  large volume increase upon oxidation of the bulk material due  to the formation of condensed oxides,  solid ZrO2(s) and BSZ liquid, where the oxide products occupy a volume 3.2 times as great as the ZrB2-SiC substrate.9 From the microstructural and chemical compositional analysis of the blisters (deformations)  we propose that the BSZ liquid forms at the reaction interface,  i.e. between a ‘‘primary’’ oxide  (a thin outer SiO2-rich borosilicate layer and an underlying porous ZrO2(s)) and the unreacted bulk material. Here it is hypothesized that the blisters  scale  \\x0c', '3654  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 91, No. 11  Fig. 3.  (a) Optical  Surface of an oxide scale on ZrB2-15 vol% SiC composite after oxidation for 4 h at 15501C. ‘‘islands’’ of ZrO2 in a film of borosilicate glass. The darker regions of the glass are rich in SiO2, which are dark because of transparency, with small dispersed ZrO2 particles barely visible. The cloudy regions are B2O3-rich borosilicate, turbid because of liquid-liquid phase separation in the glass during cooling, (b) backscattered electron image of the same region, where ZrO2 islands appear in bright contrast, SiO2-rich glassy regions in light gray contrast, and B2O3-rich regions in dark gray contrast. The small dispersed ZrO2 particles on the SiO2 appear in bright contrast.  image in reflected light, showing  form because of the large volume increase,  induced by the for mation of  the oxides during oxidation. The large volume in crease of the formed oxides induces pressure and stresses when  the oxide scale grows leading to a rupture in the ‘‘primary’’ ox ide scale. Our hypothesis is that the BSZ liquid at the reaction  interface is then squeezed up to the surface where it starts flow ing, forming the convection cells and their features.  Figure 4(d) shows a secondary electron image of a polished  cross section of a convection cell on an oxidized ZrB2-SiC sam ple. The sample was oxidized for 4 h at 15501C in a tube-furnace in ambient air, with a heating rate of 131C/min and free cooling.  The image shows the SiO2-rich external scale, in dark gray contrast, covering a ZrO2 subscale (primary ZrO2) in light gray contrast, over the unoxidized ZrB2-SiC substrate. The ZrO2 impinging on the surface (formed by precipitation of  ‘‘island’’  ZrO2) is seen in cross section to be a vertical feature containing region of glass in the middle. The glass contains B2O3, SiO2, and some ZrO2, it is also richer in boron (B) than the glass outside of  Fig. 4.  (a)-(c) Backscattered electron images of cross sections of ZrB2-15 vol% SiC composite oxidized for 3 h, showing the built up of the BSZ liquid between the ‘‘primary’’ oxide scale (SiO2-rich top layer and an under laying ZrO2) and the bulk material (ZrB2-SiC). (d) Scanning electron microscopic image of a cross section of a convection cell on ZrB2-SiC oxidized at 15501C for 4 h, showing the inner structure of a convection cell.  \\x0c', 'November 2008  Formation of Oxide Scales on ZrB2-SiC Composites  3655  Fig. 5. (a) Backscattered electron (BSE) image of a deformation with glass inside located in the cross section of the ZrB2-15%SiC composite oxidized at 15501C for 3 h; (b)-(e) the same area as in (a) imaged by electron microprobe analysis in oxygen Ka X-rays (b), zirconium La X-rays (c), silicon Ka  X-rays (d), and boron Ka X-rays (e). The scale bars on the elemental maps represent the intensity of the corresponding element. The elemental maps (b)-  (e) indicate SiO2-rich surface layer and underlying ZrO2 layer as well as the composition of the BSZ glass inside the deformation.  A  (a)  (b)  B  100 µm  Inside the “blister” (BSZ)  cps  15000  10000  5000  0  0  10       20       30        40      50         60       70       80        90       100     110     120      130    140     150     160      170      180    190     200      210       O  µm  2000  1000 0  10000  5000 0 0  2000  1000 0 0  A  0  10       20       30        40      50         60       70       80        90       100     110     120      130    140     150     160      170      180    190     200      210       10       20       30        40      50         60       70       80        90       100     110     120      130    140     150     160      170      180    190     200      210       Si  µm  Zr  µm  10       20       30        40      50         60       70       80        90       100     110     120      130    140     150     160      170      180    190     200      210       B  B  µm  Fig. 6.  (a) Backscattered electron image of a cross section of a cell on surface of a ZrB2-15 vol% SiC composite oxidized at 15501C for 3 h. The white line through the ‘‘blister’’ (deformation) indicates where the EPMA line analyses were done; the letter A indicates the start of the line scan and B the end (b) graphs of the recorded intensity ((Cps) counts per second) versus distance (mm) of the line scan.  \\x0c', '3656  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 91, No. 11  Fig. 7.  Schematic showing different planes of cross  sections possible  through a convection cell, showing the inner structure of the cell.  the ZrO2 ‘‘island’’, suggesting this to be the glass of the B2O3rich BSZ liquid. Note that the interface between the ‘‘primary’’  ZrO2 subscale and the unoxidized ZrB2-SiC substrate extends about 100 mm beneath the surface under the ZrO2 ‘‘island’’ (the center of the convection cell) but only about 50 mm away from  the center of the convection cell. Apparently this increased sub scale recession of the ZrB2-SiC substrate denotes faster oxidation at this location.  When interpreting the cross-sectional  images of the deforma tions shown in Fig. 4, one needs to consider that because the  location of the plane of polish is not known, it is hard to infer if  the features seen differ because of their location, or if they differ  because they have not developed (immature or  ‘‘before erup tion’’) or have stopped operating (‘‘extinct’’). Figure 7 is an il lustration of a convection cell  intersected by several planes of  cross section for polish. Line ‘‘a’’ in Fig. 7 intersects the center of  a BSZ liquid pipe. This can create an image similar to Fig. 4(d).  If  the plane of polish intersects  the  side of a convection cell  (through a ZrO2 ‘‘island’’ filled with liquid), i.e. line ‘‘b’’, it could create an image similar to Fig. 4(c), while a plane of polish more  remote from the pipe (line ‘‘c’’) could create an image similar to  Fig. 4(b).  IV.  Comparison with Literature  Microstructural  features closely resembling the convection cells  have been reported earlier in the literature but not interpreted as  of significance. Figure 8 compares previously published images of specimens oxidized at CNR-ISTEC by Monteverde12 with a spec imen oxidized at University of Michigan (UM). Cross sections of  HfB2-SiC-HfN and ZrB2-HfB2-SiC-HfN composites oxidized at 14501C for 20 h by Monteverde12 at CNR-ISTE are shown in  Figs. 8(a) and (b). The images show vertical ZrO2 features on top of an enhanced oxidation zone (increased thickness of ZrO2 layer). Figure 8(c) shows a ZrB2-15 vol% SiC composite fabricated at CNR-ISTEC and oxidized at 15501C for 2 h at UM.  The cross sections shown in Figs. 8(a) and (b) closely resemble  the cross section of the convection cell shown in Fig. 8(c). Note  the similarity in the morphology of  the vertical ZrO2 feature of the specimens oxidized at CNR-ISTEC to the morphology of the  ZrO2 ‘‘island’’ of arrow in Fig. 8). Also, all three specimens shown in Fig. 8 have  the specimen oxidized at UM (indicated by a  thicker ZrO2 layer (enhanced oxidation zone) under the vertical ZrO2 features. The similarity of these images indicate that the convection cells are seen in other boride-SiC materials such as  these Hf(Zr)B2-SiC-HfN composites, which have a different relative amount of in Hf(Zr)O2, SiO2, and B2O3 after oxidation. Figure 9(a) shows similar features for a ZrB2-30 vol% SiC composite oxidized for 30 min at 14001C by Rezaie et al.13 at the  University of Missouri-Rolla (UMR) previously reported in the  literature. Figure 9(b) shows a cross section of a ZrB2-15 vol% SiC composite (fabricated at CNR-ISTEC) oxidized for 1 h at 15501C at UM. The cross sections have very similar microstruc tural features: enhanced oxidation zone (thicker ZrO2 layer) under a vertical ZrO2 feature. This indicates that convection cells were formed on the ZrB2-30 vol% SiC composite during  Fig. 8.  Scanning electron microscopic images of  the cross sections of  diboride/silicon carbide composites. (a) HfB2-SiC-HfN and (b) ZrB2- HfB2-SiC-HfN composites oxidized at 14501C for 20 h by Monteverde12; (c) a ZrB2-15%SiC composite oxidized at 15501C for 2 h.  Fig. 9. Cross sections of ZrB2-SiC composites. (a) ZrB2-15%SiC oxidized at 15501C for 1 h and (b) ZrB2-30 vol% SiC oxidized at 14001C for 30 min by Rezaie et al.13  \\x0c', 'oxidation. Figure 10(a) shows a micrograph of the surface of the ZrB2-SiC specimen oxidized for 1 h at 15501C at UM. The image shows how the convection cells are spread over the surface,  forming a pattern with small micrometer-sized ZrO2 particles located between the boundaries of the cells. Figure 10(b) shows  the flow pattern of the ZrO2 particles in more detail. No micrographs of the surfaces of the specimens oxidized at CNR-ISTEC  or UMR were reported; thus comparison of these surfaces to the  UM specimens could not be done.  The  features  shown  in  previously  published micrographs  from CNR-ISTEC and UMR are suggested here to be in fact  convection cells. These findings indicate that the convection cells  do exist  for other oxidized boride-SiC materials but have not  been brought  to attention in the literature or interpreted as of  significance for the oxidation behavior of these materials.  V.  Subscale Recession: Enhanced Oxidation Regions  Now let us direct our attention to the local regions of enhances  oxidation shown in Figs. 4(d), 8 and 9. These are the areas with  more  diboride  recession  (deeper)  and  thicker  oxide  scales  (thicker ZrO2 do these form? Our hypothesis is based on inward diffusion of  layer) under  the  convection cell  features. Why  oxygen. Let us assume that  the rate of diboride oxidation is  limited by inward oxygen transport as has been reported previously in the UHTC literature.2,13,14 Areas of greater recession  demand greater inward oxygen diffusion to leave a thicker scale.  The driving force is similar; hence,  the oxygen diffusivity must  be locally higher. Here we can invoke the Stokes-Einstein relation15 between diffusivity and viscosity, derived from studies of  the Brownian motion of a solid sphere suspended in a fluid,  where the particle’s diffusivity is inversely proportional i.e. DB1/Z, more specifically15-17  to the  fluid viscosity,  D ¼ kBT  6pZr  (1)  where D is the diffusion coefficient, kB the Boltzmann constant, r is the radius of the slowest particle moving through the fluid  (the hydrodynamic  radius), and T the absolute  temperature.  Another similar equation based on the theory of absolute reaction rates by Eyring18 is sometimes preferred for silicate glasses  to relate melt viscosities to the diffusion coefficients of oxygen in molten silicates16,18,19:  D ¼ kBT  lZ  (2)  where l is \\x002). (O of 15001C is on the order of 1011 Poise (1010 Pa \\x01 s).20 With the In the SiO2-rich scale, the viscosity at oxidation temperatures  the mean jump distance of  the diffusing particle  presumed composition of BSZ liquid,10 we have estimated based  on limited data in the literature that the viscosity will be about 104 Poise (103 Pa \\x01 s)8,9,20; this is a large viscosity difference, by about factor of 107. This implies that in the local regions of the  BSZ liquid pipes, the diffusivity for inward diffusion of oxygen must be larger by a similar factor, about 107 times faster oxygen  transport than in the SiO2 scale remote from the pipe. When the convection cell forms (the deformation erupts) outward flow of  material of low viscosity (BSZ liquid) occurs and a localized in ward path for oxygen will be created. A synergy between oxygen  transport in and liquid transport out occurs. Thus we expect a  complex coupling of these phenomena.  Oxygen can diffuse through amorphous SiO2 via two mechanisms: (1) so-called network oxygen ions can diffuse through the  SiO4 tetrahedral network and (2) interstitial lar) oxygen can diffuse through the free volume of the silicate structure.21 For the network oxygen ion diffusion the Eyring can be used to estimate the oxygen diffusion coefficient by using l as  (nominally molecu the mean jump distance of the diffusing particle (O  \\x002).16,18,19 The  diffusion coefficient of  interstitial oxygen diffusion (molecular)  can be  estimated by using the Stokes-Einstein relation with r  (hydrodynamic radius) substituted by the O2 bond length (1.21 A˚ ).22 The two mechanisms can operate simultaneously and most  likely have different temperature dependencies. Thus oxidation at different temperatures could be governed by either mechanism.21  The question of whether oxygen diffuses through SiO2 as a molecule or ionically or perhaps both remains unclear, which has led to a broad range of reported diffusion coefficients.16,17,19,21-28  Figure 11 shows the calculated diffusion coefficient of oxygen  in a borosilicate melt versus  the viscosity, calculated with the  Stokes-Einstein  relation  and  the Eyring  equation. For  the  Stokes-Einstein relation the T 5 15501C,  following parameter were used;  and r 5 1.21 nm (where  the O2 bond length is used for the hydrodynamic radius of O2 22) and viscosity values are estimated by relationship extrapolated from data by Jabra and colleagues.9,20 The same viscosity values were used for the Eyring equation as well as T 5 15501C, and l 5 0.159 nm, which in SiO2 glass.19 The diffusion coefficients  is  the Si-O distance  Fig. 10.  Backscattered electron images of  the surface of a ZrB2-SiC composite oxidized at 15501C for 1 h shown earlier in cross-sectional view in Fig. 10(b); (a) low magnification (b) high magnification of the area inside the white box in (a).  1.0E-22 1.0E+00 1.0E+02  1.0E+04  1.0E+06  1.0E+08  1.0E+10  1.0E+12  Stokes-Einstein Eyring equation  η [Poise]  1.0E-20  1.0E-18  1.0E-16  1.0E-14  1.0E-12  1.0E-10  Oxygen Diffusion Coefficient vs. Viscosity   D  i  f f  u  s  i  n o  C  e o  f f  i  c  i  n e  t  [  m  s  ]  Fig. 11.  Diffusion coefficient of oxygen in a B2O3-SiO2 melt at 15501C calculated by using the Stokes-Einstein relation (diamonds) and the  Eyring equation (squares) versus the shear viscosity (logarithmic scale).  November 2008  Formation of Oxide Scales on ZrB2-SiC Composites  3657              \\x0c', 'calculated by the two equations differ on the order of about one  magnitude. The difference is not  large compared with the wide  range of values  reported in the literature; also these values are  comparable  to measured and calculated values  reported previ ously.16,17,19,21-29  Figure 12 shows the diffusion coefficient of oxygen in a boro silicate melt versus mol% of SiO2 estimated by using the Stokes- Einstein relation with T 5 15501C, r 5 0.121 nm, and viscosity values estimated from data by Jabra and colleagues.9,20 Figure 12  shows how the inward oxygen diffusion decreases with an increase  in SiO2 mol% in a B2O3-SiO2 melt. This is indicated by the low diffusion coefficient of oxygen for a pure SiO2 melt (100 mol% SiO2; D 5 1.1 \\x02 10 \\x0021 m2/s) compared with a B2O3-rich melt (79 mol% B2O3-21 mol% SiO2; D 5 1.7 \\x02 10\\x0014 m2/s), which is 107 times larger, as estimated from the large viscosity difference  mentioned above. The calculated diffusion coefficients of oxygen  in borosilicate melts (Fig. 12) is smaller than the diffusion coeffifor oxygen in ZrO2 (B10\\x0010 m2/s at 15001C) reported by cient Fox and Clyne28 supporting our previous assumption that the  borosilicate outer  layer  is  the oxygen diffusion barrier  (the rate  controlling factor) during oxidation at these temperatures.  The local enhanced oxidation regions under  the convection  cells makes the oxidation of diboride-SiC composites at  these  temperatures a nonuniform process at a microscopic scale. Ob servations of apparently uniform oxidation behavior reported in  the literature could result from the fine scale of the local events.  Just like an intergranular corrosion may be quite homogenous  on the macroscopic scale, it is though localized on a microscopic scale.30 Perhaps in the early state of the oxidation uniform ox ygen diffusion generates a pool of BSZ liquid. When the volume  of  the  liquid and the pressure  (by Pillings-Bedworth ratio)  reaches a critical amount, it erupts. While it is erupting, oxygen  can diffuse in by the same path; hence, it oxidizes faster under a  cell creating the enhanced oxidation zone.  VI.  Summary  Oxide  scale  features previously published in the  literature on  ZrB2-SiC and HfB2-SiC composites are suggested here to be in fact the proposed convection cells reported for the first time by the authors.9 These features are of great significance contributing  to the formation of complex oxide scales of these materials. The  flow of a B2O3-rich BSZ liquid to the surface, with the subsequent loss of B2O3 to evaporation, explains the formation of a glassy SiO2-rich layer on the surface commonly reported in the literature. Also the outward flow of the BSZ liquid creates a localized  inward path for oxygen due to lower viscosity that allows faster  oxidation under the convection cell. The formation of a convec tion cell  (eruption of  the ‘‘primary’’ scale) creates a leak in the  SiO2 barrier, but the BSZ liquid eventually replenishes the SiO2 scale, patching the leak creating a positive feedback.  Acknowledgments  We thank Dr. David Shifler of  the Office of Naval Research for supporting  the research under contract N00014-02-1-0034 and Drs. Alida Bellosi, Frederick  Monteverde, Gregory Hilmas, and William Farhenholtz for providing samples  and valuable discussions.  References  1F. Monteverde and A. Bellosi,  ‘‘The Resistance  to Oxidation of HfB2-SiC  Composite,’’ J. Eur. Ceram. Soc., 25, 1025-31 (2005). 2F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 3A. Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby,  in Dry  ‘‘Oxidation of  ZrB2-SiC Ceramics Under Atmospheric and Reentry Conditions,’’ Refract. Appl. Trans., 1 [2] 1-8 (2005). 4S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem,  ‘‘Evaluation  of Ultra-High Temperature Ceramics  for Aeropropulsion Use,’’  J. Eur. Ceram. Soc., 22, 2757-67 (2002). 5W. G. Fahrenholtz,  ‘‘Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 6M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurface: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 7A. Bongiorno, C. J. Fo¨ rst, R. K. Kalia, J. Li, J. Marschall, A. Nakano, M. M.  Opeka, I. G. Talmy, P. Vashishta, and S. Yip,  ‘‘A Perspective on Modeling Ma terial in Extreme Environments: Oxidation of Ultra-High Temperature Ceramics,’’  Mater. Res. Soc. Bull., 31, 410-8 (2006). 8P. C. Setze, A Review of the Physical and Thermodynamic Properties of Boric  Oxide. NACA-RM-E57B14. Lewis Flight Propulsion Laboratory, Cleveland, OH,  1957. 9S. N. Karlsdottir, J. W. Halloran, and C. E. Henderson, ‘‘Convection Patterns  in Liquid Oxide Films  on Zirconium Diboride-Silicon Carbide Composites  Oxidized at High Temperature,’’ J. Am. Ceram. Soc., 90 [9] 2863-7 (2007). 10S. N. Karlsdottir, J. W. Halloran, and A. N. Grundy, ‘‘Zirconia Transport by  Liquid Convection During Oxidation of Zirconium Diboride-Silicon Carbide  Composite,’’ J. Am. Ceram. Soc., 91 [1] 272-7 (2008). 11S. N. Karlsdottir, J. W. Halloran, F. Monteverde, and A. Bellosi, ‘‘Oxidation  of ZrB2-SiC: Comparison of Furnace Heated Coupons and Self-Heated Ribbon  Specimens’’;  in Proceedings of  the 31st International Conference on Ceramics and  Composites, Daytona Beach FL, January 21-26, 2007. Mechanical Properties and  Performance of Engineering Ceramics and Composites  III, Edited by E. Lara Curzio. Ceram. Trans., 28 [2] 327-336 (2007). 12F. Monteverde, ‘‘The Thermal Stability in Air of Hot Pressed Diboride Matrix  Composites  for Uses at Ultra-High Temperatures,’’ Corros. Sci., 47, 2020-33  (2005). 13A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Evolution of Structure  During the Oxidation of Zirconium Diboride-Silicon Carbide 15001C,’’ J. Eur. Ceram. Soc., 27 [6] 2495-501 (2007). 14F. Monteverde, ‘‘Beneficial Effects of an Ultra-Fine a-SiC Incorporation on  in Air  up  to  the Sinterability and Mechanical Properties of ZrB2,’’ Appl. Phys. A, 82, 329-37 (2006). 15A. Einstein, Investigations on the Theory of the Brownian Movement, reprinted  by Dover Publications, New York, 1926. 16Y. Liang, F. M. Richter, A. M. Davis, and E. B. Watson, ‘‘Diffusion in Silicate Melts: I. Self Diffusion in CaO-Al2O3-SiO2 at 15001C and 1 GPa,’’ Geochim. Cosmochim. Acta, 60 [22] 4353-67 (1996). 17E. M. Tanguep Njiokep and H. Mehrer, ‘‘Diffusion of 22Na and 45Ca in Ionic  Conduction in Two Standard Soda-Lime Glasses,’’ Solid State Ionics, 177, 2839-  44 (2006). 18H. Eyring,  ‘‘Viscosity, Plasticity, and Diffusion as Examples of Absolute Re action Rates,’’ J. Chem. Phys., 4 [4] 283-91 (1936). 19M. L. Ferreira Nascimento and E. D. Zanotto,  ‘‘Mechanisms and Dynamics  of Crystal Growth, Viscous Flow, and Self-Diffusion in Silica Glass,’’ Phys. Rev.  B, 73, 024209 (2006). 20R. Jabra, J. Phalippau, and J. Zarzicki,  ‘‘Synthesis of Binary Glass-Forming  Oxide Glasses by Hot-Pressing,’’ J. Non-Cryst. Solids, 42, 489-98 (1980). 21C. E. Ramberg and W. L. Worrell, ‘‘Oxygen Transport in Silica at High Tem peratures: Implications of Oxidation Kinetics,’’ J. Am. Ceram. Soc., 84 [11] 2607-  16 (2001). 22J. Read, K. Mutolo, M. Ervin, W. Behl, J. Wolfenstine, A. Driedger, and  D. Foster, ‘‘Oxygen Transport Properties of Organic Electrolytes and Performance  of Lithium/Oxygen Battery,’’ J. Electrochem. Soc., 150 [10] A1351-6 (2003). 23R. H. Doremus,  ‘‘Transport of Oxygen in Silicate Glasses,’’ J. Non-Cryst.  Solids, 349, 242-7 (2004). 24Y. Zhang, E. M. Stolper, and G. J. Wasserburg, ‘‘Diffusion of a Multi-Species  Component and its Role  in Oxygen and Water Transport  in Silicates,’’ Earth  Planet Sci. Lett., 103, 228-40 (1991). 25E. L. Williams,  ‘‘Diffusion of Oxygen in Fused Silica,’’ J. Am. Ceram. Soc.,  48 [4] 190-4 (1965). 26D. Tinker, C. E. Lesher, and I. D. Hutcheon,  ‘‘Self-Diffusion of Si and O in  Diopside Anorthite Melt at High Pressure,’’ Geochim. Cosmochim. Acta, 67 [1]  133-42 (2003). 27F. J. Norton, ‘‘Permeation of Gaseous Oxygen Through Vitreous Silica,’’ Na ture, 191, 701 (1961). 28A. C. Fox and T. W. Clyne, ‘‘Oxygen Transport by Gas Permeation Through  the Zirconia layer  in Plasma Sprayed Thermal Barrier Coatings,’’ Surf. Coat.  Technol., 184, 311-21 (2004). 29J. D. Cawley, J. W. Halloran, and A. R. Cooper, ‘‘Oxygen-18 Tracer Study of  the Passive Thermal Oxidation of Silicon,’’ Oxid. Met., 28 [1-2] 1-15 (1987). 30R. Telle, L. S. Sigl, and K. Takagi,  ‘‘Transition Metal Boride Ceramics’’; pp.  140-54 in Handbook of Ceramic Hard Materials, Vol. 1, Edited by R. Reidel.  Wiley-VCH, Weinheim, Germany, 2000.  &  0  20  40  60  80  100  Oxygen Diffusion Coefficient vs. SiO  mol.%  1.0E-21  1.0E-19  1.0E-17  1.0E-15  1.0E-13  1.0E-11  D  i  f f  u  s  i  n o  C  e o  f f  i  c  i  n e  t  [  m  s  ]  mol.% SiO     Fig. 12.  Diffusion coefficient of oxygen in a B2O3-SiO2 melt versus SiO2 mol% calculated using the Stokes-Einstein relation.  3658  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 91, No. 11                  \\x0c']"
},{
  "_id": 79,
  "PDF": "HfB2-SiC (10–20 vol _) ceramic materials Manufacture and behavior under long-term exposure to dissociated air streams.pdf",
  "Text": "['ISSN 0036(cid:2)0236, Russian Journal of Inorganic Chemistry, 2014, Vol. 59, No. 12, pp. 1361-1382. © Pleiades Publishing, Ltd., 2014. Original Russian Text © V.G. Sevastyanov, E.P. Simonenko, A.N. Gordeev, N.P. Simonenko, A.F. Kolesnikov, E.K. Papynov, O.O. Shichalin, V.A. Avramenko, N.T. Kuznetsov, 2014, published in Zhurnal Neorganicheskoi Khimii, 2014, Vol. 59, No. 12, pp. 1611-1632.  SYNTHESIS AND PROPERTIES  OF INORGANIC COMPOUNDS  HfB2(cid:2)SiC (10-20 vol %) Ceramic Materials: Manufacture  and Behavior under Long(cid:2)Term Exposure to Dissociated Air Streams  V. G. Sevastyanova, E. P. Simonenkoa, b, A. N. Gordeevc, N. P. Simonenkoa, A. F. Kolesnikovc, E. K. Papynovd, e, O. O. Shichalind, e, V. A. Avramenkod, e, and N. T. Kuznetsova  a Kurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences,  Leninskii pr. 31, Moscow, 119991 Russia  e(cid:2)mail: ep_simonenko@mail.ru b Lomonosov Moscow State University of Fine Chemical Technologies,  pr. Vernadskogo 86, Moscow, 119571, Russia c Ishlinskii Institute for Problems in Mechanics, Russian Academy of Sciences,  pr. Vernadskogo 101, block 1, Moscow, 119526 Russia d Institute of Chemistry, Far(cid:2)East Branch, Russian Academy of Sciences,  pr. Stoletiya Vladivostoka 159, Vladivostok, 690022, Russia e Far(cid:2)Eastern Federal University, ul. Sukhanova 8, Vladivostok, 690950 Russia  Received June 14, 2014  Abstract—Ultra(cid:2)high(cid:2)temperature composite materials HfB2(cid:2)SiC containing 10, 15, and 20 vol % SiC were prepared by spark plasma sintering. The behavior of the samples prepared under long(cid:2)term exposure to sub(cid:2) sonic dissociated airstreams of a high(cid:2)frequency induction plasmatron was studied. The total test time per sample was 35-42 min. Under certain exposure conditions (which were dependent on the composition of a sample), some regions of the sample were found to experience a rapid increase in temperature up to 2700°C. These regions enlarged over time, so that most of the surface area of the sample experienced exposure to tem(cid:2) peratures of up to 2500-2700°C for 19-38 min, while the rest of the surface had a temperature of up to 1800- 1900°C during almost the entire duration of the experiment. The joint use of optical microscopy, scanning electron microscopy (with EDX analysis), and X(cid:2)ray powder diffraction enabled us to study the microstruc(cid:2) ture and composition of a structurally complex oxidized layer.  DOI: 10.1134/S0036023614120250  Ceramic  composites  based  on  zirconium  or hafnium diborides and silicon carbide (ZrB2(cid:2)SiC and HfB2(cid:2)SiC, in particular when doped with other com(cid:2) ponents) are advanced by many researchers as promis(cid:2) ing materials for use in the design of thermally loaded parts, in particular, nose cones and sharp edges of wings in hypersonic aircrafts [1-7]. Due to the lucky combination of properties such as high melting tem(cid:2) peratures and nonexistence of phase transitions, resis(cid:2) tance to oxidation in air due to the ability to form a borosilicate glass barrier layer, and high heat conduc(cid:2) tivity (in particular at high temperatures), the afore(cid:2) mentioned ceramic materials are capable of with(cid:2) standing heating to temperatures higher than 2000°C under exposure to dissociated air streams [6-10]. The dopant silicon carbide plays an important role in these materials, and this implies that the silicon carbide per(cid:2) centage and distribution in the material are also of great importance in the context of oxidation stability and functional characteristics on the whole. The most recommended composites comprise 10 to 30 vol % sil(cid:2) icon carbide [11-23], although materials with higher silicon carbide percentages (of up to 45 vol % SiC) have recently been reported to be efficacious, in par(cid:2)  ticular under exposure to high(cid:2)enthalpy air flows [6, 7, 24, 25]. The most popular methods for manufacturing such materials are hot molding and spark plasma sin(cid:2) tering at temperatures of 1900-2200°C; these meth(cid:2) ods make it possible to attain 98-100% densities of the calculated values and  (as our colleagues  suggest) would thereby improve the mechanical properties and oxidative resistance due to hindered oxygen diffusion. However, our earlier studies [6, 7] showed that rela(cid:2) tively porous materials (where the calculated porosity was 19-30%) withstood long(cid:2)term exposure to disso(cid:2) ciated air flows at temperatures higher than 2000°C (up to 2700°C). We should mention that those studies were performed on a sample with a high silicon carbide percentage (25-45 vol %).  The goal of this study was to prepare HfB2(cid:2)SiC composite materials containing 10, 15, and 20 vol % silicon carbide and having relatively high porosities using spark plasma sintering and to study their behav(cid:2) ior under long(cid:2)term exposure to dissociated air flows at temperatures above 2000°C.  1361  \\x0c', '1362  SEVASTYANOV et al.  Table 1. Selected properties of HfB2(cid:2)SiC ceramic samples  Set  SiC percentage,  vol %  Density, g/cm3  Porosity*,  %  Surface roughness parameters, μm  Ra**  Ry**  1   (HfB2(cid:2)10SiC)  2 (HfB2(cid:2)15SiC)  3  (HfB2(cid:2)20SiC)  10  15  20  6.7 ± 0.1  35.6 ± 1.0  1.7 ± 0.4  4.9 ± 1.1  6.3 ± 0.1  39.7 ± 1.5  1.8 ± 0.4  5.0 ± 1.2  6.8 ± 0.1  34.9 ± 1.0  1.3 ± 0.2  3.5 ± 0.4    * Determined compared to the additively calculated density values (the HfB2 density is set equal to 10.5 g/cm3 [27] and the SiC density to 3.2 g/cm3 [28]).  ** Ra is the arithmetic mean deviation of the profile; Ry is the maximal height of the profile as determined on the baseline length of 1.25 mm.  EXPERIMENTAL  The reagents used were hafnium diboride (pure grade; particle size: 2-3 µm; aggregate size: ~20-60 µm) and silicon carbide (pure grade; average particle size: 100 µm).  Samples were manufactured on an SPS(cid:2)515S Spark Plasma Sintering System  (from Dr.Sinter(cid:2) LABTM) as follows: a premicronized mixture of HfB2 and SiC powders (where SiC percentages were 10, 15, and 20 vol % for sets 1, 2, and 3, respectively) was placed into a graphite die, compacted, evacuated, and then sintered at a temperature below 1500°C under pressure and under exposure to electric pulses with an exposure time at the maximal temperature of 20 min. Cylinder(cid:2)shaped samples (15 mm in diameter, ~5 mm high, and ~5 g in weight) were obtained in this way and were then polished.  Surface roughness parameters were determined using a TR200 (Time Group Inc.) portable roughness tester with a baseline length of 1.25 mm.  X(cid:2)ray powder diffraction studies were carried out on a Bruker D8 Advance X(cid:2)ray diffraction diffracto(cid:2) meter (CuКα radiation, 0.02° resolution).  IR transmission spectra were recorded as Nujol mulls in KBr plates on an FT(cid:2)08 Infralum IR spec(cid:2) trometer.  Scanning electron microscopy (SEM) studies were performed on an NVision 40 (Carl Zeiss) triple(cid:2)beam workstation; elemental microanalysis was carried out using an EDX Oxford Instruments energy(cid:2)dispersive attachment.  Experiments where the sample surface was exposed to a subsonic stream of dissociated air were performed on a 100(cid:2)kW VGU(cid:2)4 high(cid:2)frequency induction plas(cid:2) matron [26], in the Institute for Problems in Mechan(cid:2)  ics, with an anode supply power of 45 to 72 kW and a pressure of 100 to 300 hPa. The surface temperature was measured with a Mikron M(cid:2)770S pyrometer in the spectral ratio pyrometer mode (temperature range: 1000-3000°C; measurement spot size: ~5 mm). Tem(cid:2) perature distribution over the front surface of the sam(cid:2) ple was determined using a Tandem VS(cid:2)415U thermal imager.  RESULTS AND DISCUSSION  Characterization of Manufactured HfB2(cid:2)SiC Composite Materials  HfB2(cid:2)SiC ceramic samples containing 10, 15, and 20 vol % silicon carbide were manufactured by spark plasma sintering. Their apparent densities, calculated porosities, and roughness parameters are given  in Table 1.  Noteworthy, the HfB2(cid:2)SiC ceramic samples con(cid:2) taining high silicon carbide percentages (25-45 vol %) prepared earlier by a similar method had porosities of 20-32% [6, 7]. It follows that, provided similar man(cid:2) ufacture parameters, a reduction in SiC percentage leads to higher porosities of the resulting samples, although 10 vol % silicon carbide samples poorly fit this trend because their porosities are close to those of samples containing 20 vol % SiC. Surface roughness measurements showed values close to those recom(cid:2) mended in the literature, which are less than 1-2 µm as. The arithmetic mean deviation of the profile Ra determined on the baseline length of 1.25 mm was 1.7 ± 0.4 µm for samples of set 1 (with the minimal SiC percentage), and the maximal height of the profile Ry was ~5 µm, which  slightly exceeds  the  required parameters. The surface roughness was reduced sys(cid:2)  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1363  (а)  (b)  (c)  2000 μm (d)  1000 μm  2000 μm  (e)  1000 μm  2000 μm  (f)  1000 μm  Fig. 1. (a, b, c) External appearance and (d, e, f) surface microstructure of HfB2(cid:2)SiC ceramic samples as probed by optical microscopy. SiC percentage, vol %: (a, d) 10, (b, e) 15, and (c, f) 20.  tematically as the silicon carbide percentage increased (sets 2 and 3; see Table 1).  The external appearance of the samples was alike: they were gray cylinders having small prominencies (~20-60 µm in diameter) on their surfaces. Surface microstructure was studied by optical and scanning electron microscopy (Figs. 1, 2). As probed by SEM (Fig. 2), the surface (which consists, in all samples, mainly of well(cid:2)defined particles with sizes of 2-6 µm) has inclusions of various phase compositions com(cid:2) posed of finer particles. As the SiC percentage in ceramic samples increases, the number of these inclu(cid:2) sions (which are likely to consist mainly of silicon car(cid:2)  bide) increases, too. One can see from micrographs that the samples are rather porous.  The IR transmission spectra for all manufactured samples (as the spectra of the precursor SiC powder, too) feature, along with the absorption band ν(Si-C) at 800-850 cm-1, a low(cid:2)intensity broad absorption band with a peak at 1070-1080 cm-1 associated with the stretching vibrations ν(Si-O) of minor silicon oxide impurity on the surface of SiC particles.  The X(cid:2)ray diffraction patterns of products feature reflections from a hafnium diboride phase; low(cid:2)inten(cid:2) sity broad reflections from silicon carbide are hardly noticeable on the background, but their intensities increase in response to increasing SiC percentage.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1364  SEVASTYANOV et al.  2 μm  2 μm  2 μm  (а)  (b)  (c)  10 μm  (d)  10 μm  10 μm  (e)  10 μm  10 μm  (f)  10 μm  Fig. 2. Surface microstructure of HfB2(cid:2)SiC ceramic samples as probed by scanning electron microscopy. SiC percentage, vol %: (a, d) 10, (b, e) 15, and (c, f) 20; (d, e, f) in the phase contrast mode.  Behavior of HfB2(cid:2)SiC (10, 15, and 20 vol %) Composite  Materials under Heating with a Dissociated Air Stream  To study the behavior of the manufactured HfB2(cid:2) SiC ceramic composites under heating by a subsonic stream of dissociated air using a VGU(cid:2)4 induction plasmatron, a test sample was placed into a water(cid:2) cooled copper model whose shape was identical to the ESA standard model (which is a flat(cid:2)end cylinder with a diameter of 50 mm and a rounding edge radius of 11.5 mm). The test sample was mounted at the critical point of the water(cid:2)cooled copper model by means of  bundles of SiC whiskers so that to keep the sample from contact with the model. In order to reduce heat dissipation to the model, the sample was mounted with a 1.5(cid:2)mm protrusion from the front surface, except for sample 10V(cid:2)1, whose surface was flush with the holder surface. All experiments employed a sub(cid:2) sonic nozzle with an exit cross(cid:2)sectional diameter of 30 mm; the distance from the nozzle to the sample was also 30 mm, and the initial pressure in the high(cid:2)pressure chamber of the plasmatron was set at a level of 100 hPa. The parameters of these experiments and surface tem(cid:2) peratures are compiled in Table 2.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1365  Table 2. Results obtained from studies of the behavior of samples under exposure to subsonic dissociated air streams of a VGU(cid:2)4 high(cid:2)frequency induction plasmatron  Sample  no.  Pressure, hPa  Anodic supply  power of the  plasmatron, kW  Maximal  surface  temperature*,  °C  Overall time  of experiment,  min  Exposure time  at temperatures  above 2 000°C*,  min  Weight  change, %  Set 1 (10 vol % SiC)  10V(cid:2)11  110→120→130→140→  150→160→170→200  45→53→64  2670  10V(cid:2)2  110→120→130→140→  150→160**→150  45→53→64  2720  Set 2 (15 vol % SiC)  15V(cid:2)1  110→120→130→140  45→53→64  15V(cid:2)2  110→170  64→72→64  2740  2740  Set 3 (20 vol % SiC)  20V(cid:2)1  110→120Æ150**→130  45→53→64  2600  20V(cid:2)2  110→120→130→  170**→130  45→53→64  2540,  2600**   37  40  42  40  35  40  19  32  34  38  33  36  +1.0  +0.15  -2.0  -0.1  -3.4  -3.6  1 The sample is mounted flush with the holder and, at the first step at a temperature of ~1500°C, experienced an extreme cooling with a water stream which gave rise to a crack in the central portion of the sample, but during the second launch of the plasmatron without disassembling the system successfully withstood for more than 36 min without further cracking.    * The temperature as read from a Mikron M(cid:2)770S pyrometer.  ** For a short time.  Sample 10V(cid:2)1 of set 1. The schedule of exposure of this sample to a dissociated air stream differed appre(cid:2) ciably from that employed for the other samples: at the first step,  to  the sample which was preheated  to ~1500°C and mounted flush with the surface, cold water was admitted because of the failure of the cool(cid:2) ing system of the holder, after which cooling was stopped. The sample acquired a crack on its surface as a result, but was nonetheless forwarded to the second step of plasma chemical exposure. The variation of the surface temperature of the test sample measured with an M(cid:2)770S Mikron pyrometer (the value averaged over the surface area of ~5 mm in diameter) is shown in Fig. 3. One can see that the mean surface tempera(cid:2) ture increases slightly (by 50-100°C) as a result of a stepwise increase in anodic supply power of the plas(cid:2) matron at a fixed chamber pressure (100 hPa), as well as a subsequent increase in pressure. As probed by a Tan(cid:2) dem VS(cid:2)415U thermal imager (Fig. 4a), the surface was heated comparatively uniformly at the first steps. By  the end of the 18th minute under a pressure of 200 hPa, a local overheated region having a temperature of 1900- 1950°C appeared at the edge of the sample (Fig. 4b); this region was progressively heated to ~2700°C and grew in size, and this was responsible for a systematic increase in mean temperature (Fig. 3). Noteworthy, the hottest region has not expanded in 37 min so that to occupy the entire surface area of the sample: Fig. 4c displays the thermal image of the surface taken 1 s before the heating was switched off, and so the tem(cid:2) perature variation curve in Fig. 3 does not attain a plateau.  Figure 4d shows the external appearance of sample 10V(cid:2)1 which was withdrawn from the holder after long(cid:2)term exposure to a dissociated air stream; one can recognize some surface regions that experienced either extremely high temperatures (2600-2700°C), or relatively low temperatures (~1650-1800°C). In addition, a crack caused by abrupt water cooling at the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1366  SEVASTYANOV et al.  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  2600  2400  2200  2000  1800  1600  1400  1200  1000  0  Anodic supply power of the plasmatron  W  k  5 4  W  k  3 5  W  k  4 6  S urface te m perature  Pressure  5  10  15  20  25  30  35  Time, min  300  280  260  240  220  200  180  160  140  120  100  a  P  h  ,  e  r  u  s s  e  r  P  Fig. 3. Surface temperature of sample 10V(cid:2)1, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows (reheating).  (а)  ~16 min  (c)  37 min  1900  °С  1700  1500  1300  (b)  18 min  (d)  1700  °С  1500  1300  1100  2700  °С  2400  2200  2000  1800  1600  Fig. 4. (a-c) Thermal images of the surface of sample 10V(cid:2)1 taken in different time periods of the tests and (d) the external appearance of the sample after plasma chemical exposure.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014            \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1367  2600  2400  2200  2000  1800  1600  1400  1200  1000  Anodic supply power of the plasmatron  Surface temperature  W  k  5 4  W  k  3 5  W  k  4 6  Pressure  0  5  10  15  20  25  30  35  40  Time, min  180  170  160  150  140  130  120  110  100  a  P  h  ,  e  r  u  s s  e  r  P  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  Fig. 5. Surface temperature of sample 10V(cid:2)2, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows.  first step is clearly seen. The weight gain of the sample was 1.0%.  In general, we have to mention that temperature variation in the course of the experiment caused by changing parameters was less noticeable than usual, which can be due to some heat removal as a result of interaction with the water(cid:2)cooled holder where the sample was mounted flush with the surface of the model.  Sample 10V(cid:2)2 of set 1. Sample 10V(cid:2)2, unlike sam(cid:2) ple 10V(cid:2)1 of the same composition, was mounted in the holder so that it protruded from the holder by 1.5 mm, and so it had an appreciably different surface temper(cid:2) ature variation depending on parameters, namely on the anode supply power and pressure in the high(cid:2)pres(cid:2) sure chamber of the plasmatron (Fig. 5). The temper(cid:2) ature variation curve  features a  temperature peak (which is nonexistent in pre(cid:2)oxidized sample 10V(cid:2)1); this peak is probably associated with heat evolution in surface oxidation reactions and is leveled out by the end of the 1st minute of exposure. The mean temper(cid:2) ature experiences an insignificant increment as the power and pressure increase; with the maximal power and a pressure of 120 hPa, it begins to rise systemati(cid:2) cally while only weakly responding to a subsequent increase in pressure. On the 8th minute, the mean surface temperature exceeded 2000°C, and on the 15th minute it was stabilized at a value of 2690-2720°C; as a result, sample 10V(cid:2)2 was exposed to a dissociated air flow at a mean temperature of less than 2000°C for 32 min, of which 25 minutes were at a surface temperature of  ~2700°C.  Thermal imaging gives an explanation to the tem(cid:2) perature rise onset on the 8th minute: one can see from Fig. 6b that a progressively expanding crack appears at the edge of the sample with a temperature higher than 2000°C (on the 9th minute, higher than 2600°C). An enlargement of  the surface area of  this region  to occupy the entire surface of sample 10V(cid:2)2 at the end of the 15th minute (Fig. 6), is responsible for the increase of the mean temperature detected by the pyrometer (Fig. 5).  The transient increase in pressure in the high(cid:2)pres(cid:2) sure chamber of the plasmatron from 150 to 160 GPa in no way influenced the surface temperature, which may serve as indirect evidence that the surface was strongly catalytic.  Figure 6d shows that the surface of sample 10V(cid:2)2 after testing had a uniform white color of oxidation products of HfB2(cid:4)SiС composite. The weight gain of the sample after 40 min of exposure was 0.15%.  Sample 15V(cid:2)1 of set 2. This sample was subjected to the longest exposure: the overall exposure time was 42 min (Fig. 7).  We must mention that the temperature schedule of sample 15V(cid:2)1 at the initial stage resembles that for sample 10V(cid:2)2: an  insignificant  increase  in surface temperature occurred with a stepwise rise in anodic supply power of the plasmatron. The mean tempera(cid:2) ture started to increase on the 5th-6th minute, which was due to the appearance and enlargement of local overheated areas whose  temperature  far exceeded 1900-2000°C (Fig. 8a). This processes become most active starting at the 7th-8th minute after the pressure  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014            \\x0c', '1368  (а)  2 min  (c)  15 min  2000 °С  1800  1600  1400  1200  SEVASTYANOV et al.  (b)  8 min  1600  °С  1400  1200  1000  0280  °С  2500  0  2000  1500  (d)  Fig. 6. (a-c) Thermal images of the surface of sample 10V(cid:2)2 taken in different time periods of the tests and (d) the external appear(cid:2) ance of the sample after plasma chemical exposure.  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  Anodic supply power of the plasmatron  Surface temperature  Pressure  W  k  5 4  W  k  3 5  W  k  4 6  2600  2400  2200  2000  1800  1600  1400  1200  1000  0  5  10  15  20  25  30  35  40  Time, min.  200  150  100  50  0  a  P  h  ,  e  r  u  s s  e  r  P  Fig. 7. Surface temperature of sample 15V(cid:2)1, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows.  in  the high(cid:2)pressure chamber of  the plasmatron increased from 110 to 120 hPa (Fig. 8b); at the end of the 13th minute, the surface temperature of the sample equalized over all areas at ~2640-2740°C (Fig. 8c). The  highest pressure in the high(cid:2)pressure chamber of the plasmatron in this case was 140 hPa (Figs. 8d, 8e).  After sharp cooling as a result of switched(cid:2)off heat(cid:2) ing, the surface temperature of sample 15V(cid:2)1 lowered  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014            \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1369  (а)  (c)  5 min  13 min  8 min  40 min  (b)  (d)  2000  °С  1800  1600  1400  2700  °С  2400  2100  1800  2600  °С  2300  2000  1700  2600  °С  2300  2000  1700  (e)  42 min  2500  (f)  °С  2000  1600  Fig. 8. (a-e) Thermal images of the surface of sample 15V(cid:2)1 taken in different time periods of the tests and (f) the external appear(cid:2) ance of the sample after plasma chemical exposure.  from ~2640-2740 to 1050-1100°C in 5 s, which did not bring about the disintegration or bundling of the sample. The external appearance of sample 15V(cid:2)1 in the holder after exposure to a dissociated air stream for 42 min is shown in Fig. 8a. The weight loss of the sam(cid:2) ple was 2.0%, which may be associated with intense evaporation of boron(cid:2) and silicon(cid:2)containing prod(cid:2) ucts from the surface at extremely high temperatures (2600-2740°C) and relatively low pressures in the high(cid:2)pressure chamber of the plasmatron.  Sample 15V(cid:2)2 of set 2. The plasma chemical exper(cid:2) iment involving sample 15V(cid:2)2 was purposed to study its behavior under rapid heating to high temperatures. For this purpose, the sample was introduced into a dis(cid:2) sociated air stream with the 64(cid:2)kW anodic supply power of the plasmatron, which was the maximal value for all of the experiments described here; the pressure in the chamber also increased from 100 to ~167-170 hPa in 1.5-2 min.  The surface temperature variation of the sample is shown in Fig. 9. One can see that the surface temper(cid:2)  ature at the 2nd-3rd minute exceeded 2000°C as a result of the rapid expansion of a local overheated area (having a temperature of ~2700°C) that emerged on the periphery of the sample whose temperature was 1700-1800°C (Figs. 10a-10c). Thus, in 11 minutes the  surface  temperature was at a  level of 2700- 2710°C. The transient (from 13th through 14 minute) increase in the anodic supply power of the plasmatron to 72 kW gave rise to a ruse in temperature to 2740- 2750°C, and after the 64(cid:2)kW power was recovered, the temperature  returned  to  the  previous  level  of ~2700°C. On the 15th minute, one can trace, in the thermal image, the appearance of a region with a slightly lower (by ~100°C) temperature, which sur(cid:2) vived to the end of the experiment. The external appear(cid:2) ance of the sample in the holder after tests is shown in Fig. 10d. The overall weight loss of the sample in 40 min of exposure to a dissociated air stream was 0.1%.  Sample 20V(cid:2)1 of set 3. Exposure of this sample to dissociated air streams was also controlled through changing the anodic supply power of the plasmatron  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1370  SEVASTYANOV et al.  Anodic supply power of the plasmatron  72 kW  Surface temperature  W  k  4 6  Pressure  W  k  4 6  200  190  180  170  160  150  140  130  120  110  100  2600  2400  2200  2000  1800  1600  1400  1200  1000  a  P  h  ,  e  r  u  s s  e  r  P  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  0  10  20 Time, min  30  40  Fig. 9. Surface temperature of sample 15V(cid:2)2, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows.  (а)  (c)  116 s  11 min  (b)  205 s  (d)  2200  °С  2000  1800  1600  1400  2700 °С  2400  2100  1800  2700  °С  2400  2100  1800  Fig. 10. (a-c) Thermal images of the surface of sample 15V(cid:2)2 taken in different time periods of the tests and (d) the external appearance of the sample after plasma chemical exposure.  from 45 to 53 and then to 64, followed by stepwise  and pressure values to exceed 2000°C in the beginning  increasing pressure in the chamber (Fig. 11). One can  of the 3rd minute (Fig. 12a), and this strongly distin(cid:2)  see that the surface temperature increases as early as at  guish the behavior of this sample from the behavior of  the first stage of the experiment with minimal power  samples of sets 1 and 2. At the end of the 2nd minute,  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014          \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1371  2600  2400  2200  2000  1800  1600  1400  1200  1000  Anodic supply power of the plasmatron  Surface temperature  W  k  4 6  W  k  5 4  W  k  3 5  Pressure  160  150  140  130  120  110  0  5  10  15  20  Time, min  25  30  100  35  a  P  h  ,  e  r  u  s s  e  r  P  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  Fig. 11. Surface temperature of sample 20V(cid:2)1, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows.  2 min 50 s  2500  2300  °С  2000  1700  (а)  (c)  1 min 20 s  35 min  2000  (b)  °С  1800  1600  1400  2600 °С  2300  2000  1700  (d)  Fig. 12. (a-c) Thermal images of the surface of sample 20V(cid:2)1 taken in different time periods of the tests and (d) the external appearance of the sample after plasma chemical exposure.  the sample was pushed by the gas flow into the holder (the protrusion size was ~0.5 mm as a result), then the surface temperature acquired a value of 2370-2440°C with a rise to 2550-2610°C, except for small areas that had lower temperatures (1800-1900°C), which were reduced in size only insignificantly as the power and pressure increased further (Figs. 12b, 12c).  We must mention that when the anodic supply power of the plasmatron increased from 53 to 64 kW, the mean temperature increased little (by ~50°C), and an appreciable increase in pressure from 120 to 153 hPa had almost no effect on the surface temperature of sample 20V(cid:2)1 (the rise was ~15°C), which may indi(cid:2) cate the high surface catalyticity in surface recombina(cid:2)  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014            \\x0c', '1372  SEVASTYANOV et al.  Anodic supply power of the plasmatron  C  °  ,  e  r  u  t  a  r  e  p  m  e  T  2600  2400  2200  2000  1800  1600  1400  1200  1000  W  k  5 4  W  k  3 5  64 kW  170  160  150  140  130  120  110  0  5  10  15  20  25  30  35  100  40  Time, min  a  P  h  ,  e  r  u  s s  e  r  P  Fig. 13. Surface temperature of sample 20V(cid:2)2, averaged over the central region ~5 mm in diameter (as measured with a Mikron M(cid:2)770S pyrometer), chamber pressure, and anodic supply power in the plasmatron during exposure to dissociated air flows.  tion of oxygen and nitrogen atoms. The external appearance of the sample in the holder after tests is shown in Fig. 12d: the larger surface area which corre(cid:2) sponds to the high(cid:2)temperature test is white in color, while low(cid:2)temperature areas have a gray tint (thermal images and micrographs were taken from opposite sides of the unit). The weight loss 35 min of exposure to a dissociated air stream was 3.4%.  Sample 20V(cid:2)2 of set 3. The exposure protocol for this sample repeats that for compositionally similar sample 20V(cid:2)1 (Fig. 13). The difference consists in that sample 20V(cid:2)2 was partially pushed into the holder by the gas flow as early as at the initial stage, resulting in an inclination of the front surface and thereby improv(cid:2) ing heat dissipation and slowing down the full heating of the surface of the sample. Thus, the attainment of temperatures higher than 2000°C and the onset of an active rise in temperature occurred at the third heating stage when the anodic supply power of the plasmatron was 64 kW and the pressure had a minimal value of 110 hPa (Fig. 14b), although a slow increase in mean tempera(cid:2) ture of the sample was observed as early as at the first stage.  One can see that the surface temperature of the sample was 2530-2550°C most of time (for more than 35 min), and an appreciable increase in pressure to ~170 hPa almost did not change this value (the rise was ~15°C). The weight loss of the sample was 3.6%.  Characterization of HfB2(cid:2)SiC (10, 15 and 20 vol % SiC)  Samples after Exposure to Dissociated Air Streams  Surface roughness was measured in all samples after plasma chemical exposure. Different regions of  sample 10V(cid:2)1, where temperatures were appreciably different during the experiment were found to have different roughnesses. In a hot area which experienced temperatures of 2600-2700°C,  the roughness was reduced: Ra was 1.1 µm, i.e., more than twice lower than for the intact sample. In the areas where the temper(cid:2) atures were ~1650-1800°C, Ra values were ~3.0 µm. For sample 10V(cid:2)2 whose surface temperature exceeded 2600°C during the tests, the parameter Ra increased twice to become also 3.0 µm. For samples of set 2, the parameter Ra also increased twofold to become 3.2 (in sample 15V(cid:2)1) and 4.7 µm (in sample 15V(cid:2)2). In sam(cid:2) ple 20V(cid:2)1, both the hot and the cold areas experienced an increase in Ra (to 3.4 and 2.9 µm, respectively), while in sample 20V(cid:2)2 (which has the same composi(cid:2) tion and exposure protocol),  the mean roughness remained almost unchanged: Ra ~1.5 µm. The X(cid:2)ray powder diffraction patterns recorded from the front surface showed that monoclinic HfO2 was the only crystalline phase identified in all samples except for sample 10V(cid:2)1 (Fig. 15).  For sample 10V(cid:2)1, additional reflections corre(cid:2) sponding  to hafnium  silicate HfSiO4 phase  are observed. Likely, from the data earlier obtained for a sample having an appreciably higher SiC percentage SiC (45 vol %) [7], one can infer that the formation of this phase is typical in case of high pressures (200 hPa) and when the surface temperature during the plasma chemical experiment is lower than 1800-1900°C in the cold areas formed for sample 10V(cid:2)1.  This  inference  is  indirectly  supported by  the absence of reflections from the HfSiO4 phase in the X(cid:2)ray diffraction pattern of compositionally similar sample 10V(cid:2)2, whose surface was uniformly heated to  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014          \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1373  4 min 40 s  2700  °С  2400  2100  1800  (а)  1 min  (c)  40 min  (b)  (d)  1800 °С  1600  1400  2700 °С  2400  2100  1800  Fig. 14. (a-c) Thermal images of the surface of sample 20V(cid:2)2 taken in different time periods of the tests and (d) the external appearance of the sample after plasma chemical exposure.  HfO2  HfSiO4  I  20V(cid:4)2  20V(cid:4)1  15V(cid:4)2  15V(cid:4)1  10V(cid:4)2  10V(cid:4)1  10  20  30  40  2θ, deg  50  60  Fig. 15. X(cid:2)ray diffraction patterns recorded from the front surfaces of HfB2(cid:2)SiC (10, 15, and 20 vol % SiC) samples after exposure to dissociated air flows.  2690-2720°C during exposure to a dissociated air stream. Further, the exposure conditions for sample 10V(cid:2)2 favored removal of boron(cid:2) and silicon(cid:2)contain(cid:2) ing oxidation products: the experiment was carried out at appreciably lower pressures in the high(cid:2)pressure chamber and longer exposure times at maximal tem(cid:2) peratures, which is also proven by the difference in weight change as a result of concurrent oxidation and  evaporation of liquid products (components of newly formed borosilicate glass).  For samples of sets 2 and 3, a hafnone phase was absent on the front surface, regardless of whether areas with temperatures lower than 2000°C were observed or no; this may be explained by longer exposure times (at least 33 min) at a mean  surface  temperature or ~2550-2740°C and at relatively low pressures (≤130-  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1374  SEVASTYANOV et al.  (а)  (b)  1000 μm  (c)  1000 μm  (e)  1000 μm  1000 μm  (d)  1000 μm  (f)  1000 μm  Fig. 16. External appearance of the front surfaces of HfB2(cid:2)SiC samples after exposure to dissociated air flows: (a) sample 10V(cid:2)1, (b) sample 10V(cid:2)2, (c) 15V(cid:2)1, (d) 15V(cid:2)2, (e) sample 20V(cid:2)1, and (f) sample 20V(cid:2)2.  140 hPa, with transient increases). Some exception is sample 15V(cid:2)2; in the experiment with this sample the pressure was 170 hPa, but the nonappearance of the HfSiO4 phase is likely to be due to a longer exposure time (38 min) at a temperature higher than 2000°C.  Figure 16 shows the external appearance of sam(cid:2) ples after exposure to dissociated air streams. For sam(cid:2) ple 10V(cid:2)1, one can clearly see the crack resulting from rapid cooling with water at the first stage of exposure to air plasma; one can also see an area that had a lower temperature (not higher  than 1800°C) during  the experiment, as distinct from the most surface area of the sample. For the other samples, the oxidized sur(cid:2) face has the appearance of white porous ceramics. Exceptions are the samples with the highest silicon carbide percentage (20 vol %), whose surface also has gray areas where temperatures were appreciably lower than over most surface areas of the samples, with a characteristic glassy luster. It is worth noting that the crack observed in sample 15V(cid:2)2 was not formed in the course of exposure; rather it appeared during storage of the sample in 2 weeks after the test (it is likely that stress  relaxation or damage occurred during  the removal from the holder).  Scanning electron microscopy was used to gain more details of the surface microstructure of samples after plasmatron experiments (Figs. 17-22).  Very interesting is the situation where the surface of a sample, apart from a greater area exposed to temper(cid:2) atures of 2600-2700°C, has some areas where the temperature during the experiment was relatively low (less than 1800-1900°C). Sample 10V(cid:2)1 experienced  the  least high(cid:2)temperature exposure:  the exposure time at a mean temperature above 2000°C for this sample was as short as 19 min; there were opportuni(cid:2) ties for more active heat exchange with the holder, because the sample was mounted in the holder without protrusion; during half of the total test time, the high(cid:2) est pressure of all the experiments (200 hPa) was main(cid:2) tained in the high(cid:2)pressure chamber of the plasma(cid:2) tron, and this should intensify oxidation processes and counteract the high(cid:2)temperature evaporation of the components of borosilicate glass from the surface. As a result one can see that the complete evaporation of boron(cid:2) and silicon(cid:2)containing components did not occur even in the central region (Figs. 17a, 17b) where temperatures of 2600-2700°C were observed, and borosilicate glass occurs in the pores of a refractory HfO2 skeleton at relatively small depths. An entirely dif(cid:2) ferent picture was observed on the surface areas where the temperature did not exceed 1800°C (Figs. 17c, 17d): their surfaces are fully covered with the glass extruded from the pores of by excessive pressure of SiO(g) and CO(g) [29-31], the surface of which bears solidified bubbles indicating the occurrence of gas formation or evaporation. Phase contrast data allow us to say that the conditions created on such surface areas kept the HfO2 skeleton from yielding to the surface (this is also indicated by EDX data: the hafnium percentage is as low as 3.6 at % against the silicon percentage of 18.5 at %). Matching X(cid:2)ray powder diffraction data with  the existing picture, we may assume that the concentra(cid:2) tion conditions (sufficiently high SiO2 percentage) and temperature conditions enabled hafnone HfSiO4  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1375  (а)  (c)  5 μm  (b)  5 μm  5 μm  (d)  5 μm  Fig. 17. Surface microstructure of sample 10V(cid:2)1 of set 1 after exposure to dissociated air flows (as probed by SEM): (a, c) surface morphology as probed by a secondary electron detector for (a) hot area and (c) cold area and (b, d) in the atomic number averaged contrast mode.  to crystallize upon rapid cooling; this is not in contrary to the HfO2-SiO2 phase diagram [32]. For compositionally similar sample 10V(cid:2)2 which also comprises 10 vol % SiC and was subjected to a more severe exposure [longer exposure times (~25 min at 2600-2700°C and 32 min at temperatures higher than 2000°C) and an appreciably lower pressure in the high(cid:2)pressure chamber of  the plasmatron, which should intensify the volatilization of boron and silicon oxides], a porous structure was formed throughout the entire surface area of the sample; this structure con(cid:2) sisted mostly of hafnium dioxide (silicon was not detected by EDX on the surface). Probably, the pro(cid:2) tective borosilicate glass layer was in the deeper lying regions of the multilayer near(cid:2)surface oxidized region of the sample as a result of active evaporation from the surface.  For sample 15V(cid:2)1 (Fig. 19), a higher porosity of surface microstructure was observed (there was a high proportion of large pores with sizes of 20-50 µm). Energy(cid:2)dispersive analysis of the 6(cid:2)mm2 central area showed that this area consisted mostly of HfO2 with an insignificant silicon impurity.  A similar situation  is for sample 15V(cid:2)2, which belongs to the same set 2 and has a very similar surface  microstructure (Fig. 20): an incipiently melted porous HfO2 skeleton was observed and no borosilicate glass was detected on the surface (EDX did not show silicon traces). It is likely that the higher pressure in the high(cid:2) pressure chamber of the plasmatron compared to that in the experiment on sample 15V(cid:2)1 (for sample 15V(cid:2)1, Р ≤ 140 hPa and a transient hypersonic conditions were created) were compensated for at Р = 170 hPa by a  long(cid:2)term exposure at  temperature higher  than 2700°C (32 min against 25 min for sample 15V(cid:2)1).  For both samples of set 3, conditions were created such that areas having temperature below 1850°C were formed on the surface, most area of which was heated  to 2500-2600°C. This  impacted both  the microstructure and chemical composition of these areas. It is pertinent in this regard that the cold areas were also exposed for 34 and 38 min at relatively low pressures in the high(cid:2)pressure chamber (mostly of 130 hPa), which should have favor the evaporation of more volatile components.  So, sample 20V(cid:2)1, whose exposure was less long than for its analogue 20V(cid:2)2, mostly formed the micro(cid:2) structure of a porous ceramic skeleton of HfO2 (as for the other samples, except for 10V(cid:2)1), which is verified by  elemental  analysis  (no  silicon  impurity was  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1376  SEVASTYANOV et al.  (а)  (c)  5 µm  (b)  5 µm  20 µm  (d)  2 µm  Fig. 18. Surface microstructure of sample 10V(cid:2)2 of set 1 after exposure to dissociated air flows (as probed by SEM): (a, c, d) sur(cid:2) face morphology as probed by a secondary electron and (b) in the atomic number averaged contrast mode.  (а)  (c)  5 µm  (b)  5 µm  20 µm (d)  2 µm  Fig. 19. Surface microstructure of sample 15V(cid:2)1 of set 2 after exposure to dissociated air flows (as probed by SEM): (a, c, d) sur(cid:2) face morphology as probed by a secondary electron and (b) in the atomic number averaged contrast mode.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1377  (а)  (c)  5 μm  (b)  5 μm  20 μm (d)  2 μm  Fig. 20. Surface microstructure of sample 15V(cid:2)2 of set 2 after exposure to dissociated air flows (as probed by SEM): (a, c, d) sur(cid:2) face morphology as probed by a secondary electron and (b) in the atomic number averaged contrast mode.  detected on ~6(cid:2)mm2 central surface area). On cold areas (Figs. 21c, 21d), one can see that the surface(cid:2) yielding HfO2 skeleton is partially filled with borosili(cid:2) cate glass. Energy(cid:2)dispersive analysis showed that, despite the absence of silicon(cid:2)containing crystalline phases, silicon oxide on the surface (Fig. 15) in this area (silicon percentage is ~19 at %) prevails over the  latter performs  the  role of a  skeleton (hafnium percentage is ~3-4 at %), which starts to manifest itself over partially evaporated and deeper buried glass.  HfO2;   The surface microstructure of sample 20V(cid:2)2 may be described in the same manner: a porous hafnium dioxide surface layer dominates (silicon as absent on the surface as probed by EDX analysis), and is accom(cid:2) panied with glass in the areas that were heated to lower temperatures during the plasma chemical experiment (Fig. 22). The following is noteworthy: since this sam(cid:2) ple experienced a longer exposure to a dissociated air stream, the degree of borosilicate glass evaporation from the surface was higher. Figures 22c and 22d show that, unlike the situation with sample 20V(cid:2)1, large pores are formed in the HfO2 skeleton and glass only slightly wets it because of having high wettability on zirconium and hafnium dioxides. Energy(cid:2)dispersive  analysis of these areas showed higher hafnium per(cid:2) centages compared to sample 20V(cid:2)1 (15-20 at %), but the silicon fraction was also considerable (6-11 at %).  On the whole, we may say that porous ceramic hafnium dioxide skeleton was formed on the surfaces of samples of all sets (containing 10, 15, and 20 vol % SiC), which were heated by dissociated air streams to temperatures of 2550-2700°C. When there were areas with appreciably  lower temperatures, the chemical composition and surface microstructure were strongly influenced by specific exposure conditions, primarily the pressure in the high(cid:2)pressure chamber of the plas(cid:2) matron (which influences the bulk oxygen concentra(cid:2) tion and the conditions of possible evaporation of boron and silicon oxides from the surface) and the exposure length and temperature.  The oxidation of HfB2(cid:2)SiC materials in the bulk was studied using optical microscopy (Fig. 23) and scan(cid:2) ning electron microscopy (Fig. 24). Figure 23 shows, as examples, sectional images of samples 10V(cid:2)2, 15V(cid:2)1, and 20V(cid:2)1 (panel c), which were exposed to dissoci(cid:2) ated air streams under similar conditions (especially in regard of the variation of anodic supply power of the plasmatron). One can see that a multilayer oxidized area was formed with a thickness ranging from about  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1378  SEVASTYANOV et al.  (а)  (c)  5 μm  (b)  5 μm  5 μm (d)  5 μm  Fig. 21. Surface microstructure of sample 20V(cid:2)1 of set 3 after exposure to dissociated air flows (as probed by SEM): (a, c) surface morphology as probed by a secondary electron detector for (a) hot area and (c) cold area and (b, d) in the atomic number averaged contrast mode.  100 µm (for areas where the temperature did not exceed 1850-1900°C in the course of plasma chemi(cid:2) cal exposure; see Fig. 23c, right(cid:2)hand image) to 800- 1000 µm for hot areas where the surface was heated to and exposed for more than 20-30 min at 2550- 2700°C. A rather porous surface layer is formed (likely based on refractory HfO2), underlain by more dense layers where the pores of the HfO2 skeleton was fully filled with oxygen diffusion protecting borosilicate glass produced by  the oxidation of  the HfB2(cid:2)SiC material.  Some  samples  have  glass(cid:2)filled  pores extended along the surface.  SEM data for sections of the same samples (Fig. 24) make it possible to have insight into more details of the microstructure and to estimate the elemental compo(cid:2) sition of each layer. The pore morphology changes considerably in going from the surface to inner regions of a sample. So (Fig. 25), on the surface dominated by hafnium dioxide, large (5-50 µm) pores prevail. At the same time, vertically extended pores occur in var(cid:2) ious areas of the oxidized layer, which likely served for gas evolution in the course of plasma chemical expo(cid:2) sure. At the boundary with the silicon carbide depleted area, there are horizontally extended pores which can  in future serve for the exfoliation of the oxidized por(cid:2) tion of the sample (for the reason that the silicon car(cid:2) bide depleted area of  the ceramic sample  is very porous).  Table 3 lists the thicknesses of oxidized areas for all of the samples studied. One can see that the behavior under similar parameters of exposure to dissociated air streams and the thickness of the oxidized area vary depending on  the composition of  the HfB2(cid:2)SiC ceramic composite sample. When a sample is arranged with a protrusion of ~1.5 mm, the anodic supply power of the plasmatron changes stepwise, and then the pressure increases, the samples having the minimal SiC percentage (10 vol %; sample 10V(cid:2)1), start rapidly cooling at a higher power (64 kW), and when the pres(cid:2) sure increases to 120-130 hPa. For samples of set 2 (15 vol % SiC: sample 15V(cid:2)1), a rapid rise in surface tem(cid:2) perature is observed under milder exposure: the power is also 64 kW and the pressure is lower (110-120 hPa). With this, the maximal surface temperature for samples of sets 1 and 2 is attained at temperatures higher than 2700°C, which is close to the melting temperature of the highest melting component of the oxidized layer (HfO2, Tm = 2780 ± 30 [33] or 2820 [34]); oxidized  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1379  (а)  (c)  5 μm  (b)  5 μm  5 μm (d)  5 μm  Fig. 22. Surface microstructure of sample 20V(cid:2)2 of set 3 after exposure to dissociated air flows (as probed by SEM): (a, c) surface morphology as probed by a secondary electron detector for (a) hot area and (c) cold area and (b, d) in the atomic number averaged contrast mode.  layers for these samples have roughly equal thick(cid:2) nesses. The samples of set 3 containing 20 vol % sili(cid:2) con carbide (sample 20V(cid:2)1 and sample 20V(cid:2)2) are dis(cid:2) tinguished by a very rapid increase in temperature under molder exposure conditions even at the least power (45 kW) and a pressure of 110 hPa. The maxi(cid:2) mal surface temperature for these samples does not exceed 2600-2610°C, which is lower than for samples of sets 1 and 2, so that both the maximal and average thicknesses of the oxidized portion are 200-300 µm smaller. Pressure in the course of the experiment is also of importance because both the oxygen concen(cid:2) tration and the activity of evaporation of volatile boron and silicon oxides are pressure dependent, and this can be manifested in the weight change of the sample after tests.  On the whole we may say that, regardless of the composition and porosity, all samples withstood expo(cid:2) sure to dissociated air streams without being disinte(cid:2) grated, despite the high temperature attained (2600- 2740°C). Rapid cooling when heating was switched off (by more than 1500°C in 3-5 s) likewise did not gave rise to cracking or bundling immediately after the tests. An HfO2  thermal barrier  surface  layer was  observed for all samples (on the area heated to 2600- 2700°C). On the areas where the temperature did not exceed 1800-1900°C, there were mixtures of HfO2 (a refractory skeleton: solid pillars) and hafnone, which  Table 3. Thicknesses l of oxidized regions (omitting SiC(cid:2) depleted layers) for HfB2(cid:2)SiC ceramic material samples af(cid:2) ter exposure to dissociated air streams  Sample no.  lmax, μm  lmin, μm  lavg, μm*  10V(cid:2)1  10V(cid:2)2  15V(cid:2)1  15V(cid:2)2  Set 1 (10 vol % SiC)  390  934  30**  480  Set 2 (15 vol % SiC)  950  980  375  340  Set 3 (20 vol % SiC)  20V(cid:2)1  20V(cid:2)2***  710  720  120**  160**      * In the hot area.    ** In the cold area.  140-180  700-900  650-900  700-900  400-650  300-600  *** The oxidized portion of sample 20V(cid:2)2 cleaved from the region depleted of silicon carbide during section preparation.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1380  SEVASTYANOV et al.  (а)  (b)  (c)  1000 μm  1000 μm  1000 μm  1000 μm  1000 μm  1000 μm  Fig. 23. Sectional microstructures of samples as probed by optical microscopy: (a) sample 10V(cid:2)2, (b) sample 15V(cid:2)1, and (c) 20V(cid:2)1.  (а)  (b)  (c)  1000 μm  1000 μm  1000 μm  Fig. 24. Sectional microstructures of samples as probed by SEM (InLens secondary electron detector): (a) sample 10V(cid:2)2, (b) sample 15V(cid:2)1, and (c) sample 20V(cid:2)1.  crystallized  from  the  low(cid:2)boron(cid:2)oxide borosilicate glass melt; that glass was extruded to the surface under an excess pressure of silicon and carbon monoxides (liquid roof) and actively evaporated to progressively open glass(cid:2)wetted pores of HfO2 solid pillars because of the high temperature of the liquid roof and a rela(cid:2) tively low pressure in the high(cid:2)pressure chamber of the plasmatron. These phenomena are also intrinsic to samples with higher silicon carbide percentages (25, 35, and 45 vol %); the results obtained by exposure of  those samples to dissociated air streams have been described earlier [6, 7].  CONCLUSIONS  We employed spark plasma sintering to manufac(cid:2) ture HfB2(cid:2)SiC  ultra(cid:2)high(cid:2)temperature  composite materials having diverse silicon carbide volume per(cid:2) centages (10, 15, and 20 vol %) and the tailored poros(cid:2) ities of ~35-40%. We have not observed significant  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', 'HfB2(cid:2)SiC (10-20 VOL %) CERAMIC MATERIALS  1381  another in the phase and elemental compositions and in surface and cross sectional microstructures.  Our studies prove the potential of HfB2(cid:2)SiC porous materials manufactured using spark plasma sintering for use under heating, in particular, under exposure to dissociated air flows at ultrahigh surface temperatures and prove the need for continuing systematic research in this field.  ACKNOWLEDGMENTS  This study was supported by the Presidential Grant for young scientists (MK(cid:2)1435.2013.3) and the Rus(cid:2) sian Foundation for Basic Research (13(cid:2)03(cid:2)12206(cid:2) ofi_m).  REFERENCES  1. E. P. Simonenko, D. V. Sevast’yanov, and N.P. Simo(cid:2) nenko et al., Russ. J. Inorg. Chem. 58, 1669 (2013).  2. M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, J. Mater. Sci. 39, 5887 (2004).  3. F. Monteverde and R. Savino, J. Am. Ceram. Soc. 95, 2282 (2012).  4. T. H. Squire and J. Marschall, J. Eur. Ceram. Soc. 30, 2239 (2010).  5. W. G. Fahrenholtz and G. E. Hilmas, Int. Mater. Rev. 57 (1), 61 (2012).  6. V. G. Sevast’yanov, E. P. Simonenko and A. N. Gordeev, et al., Russ. J. Inorg. Chem. 58, 1269 (2013).  7. V.G. Sevastyanov, E. P. Simonenko and A. N. Gordeev, et al., Russ. J. Inorg. Chem. 59, 1298 (2014).  8.  J. Marschall, D. A. Pejakovic (cid:5), W. G. Fahrenholtz, et al., J. Thermophys. Heat Transfer 26, 559 (2012).  9. M. Gasch, D. Ellerby, E. 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Sectional microstructure of sample 10V(cid:2)2 (hot area) as probed by SEM (SE2 secondary electron detector).  defects on the surfaces of samples; the arithmetic mean deviation of the profile derived from surface roughness was ~1.5-2.0 µm. The elemental and phase compositions of the materials have been determined.  HfB2(cid:2)SiC  ceramic  composite  samples  were exposed to high(cid:2)enthalpy dissociated air streams of a VGU(cid:2)4 induction plasmatron with an anode supply power of 45 to 72 kW and a pressure of 100 to 200 hPa in the plasmatron high(cid:2)pressure chamber. Under cer(cid:2) tain parameters of the experiment (which were appre(cid:2) ciably differentiated for samples of different sets), local overheating areas were observed to appear (as a rule, the temperature at the edges of the sample con(cid:2) siderably exceeded 2000°C); these areas progressively merged during exposure to dissociated air streams to occupy the entire or almost entire surface area of the sample. The temperature acquired values of 2550- 2600 (in samples containing 20 vol % SiC) or 2700- 2740°C (in samples containing 10 or 15 vol % SiC). We have shown that samples that experienced expo(cid:2) sure to temperatures of 2600-2700 and those exposed to  differed  appreciably  from  one  1800-1900°C   RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c', '1382  SEVASTYANOV et al.  21. S. Gangireddy, J. W. Halloran, and Z. N. Wing, J. Eur. Ceram. Soc. 33, 2901 (2013).  22. W. Tan, C. A. Petorak, and R. W. Trice, J. Eur. Ceram. Soc. 34, 1 (2014).  23. C. Carney, A. Paul, S. Venugopal, et al., J. Eur. Ceram. Soc. 34, 1045 (2014).  24. M. Tului, B. Giambi, S. Lionetti, et al., Surf. Coat. Technol. 207, 182 (2012).  25. P. A. Williams, R. Sakidj, J. H. Perepezko, and P. Ritt, J. Eur. Ceram. Soc. 32, 3875 (2012).  26. A. N. Gordeev and A. F. Kolesnikov, Challenging Topics of Mechanics: Physicochemical Mechanics of Liquids and Gases (Nauka, Moscow, 2010) [in Russian].  27. D. E. Wiley, W. R. Manning, and O. Hunter, Jr., J. Less(cid:2) Common Met. 18, 149 (1969).  28. T. Kawamura, Miner. J. (Japan) 4, 333 (1965).  29.  J. Li, T. J. Lenosky, C. J. Foörst, et al., J. Am. Ceram. Soc. 91, 1475 (2008).  30. T. A. Parthasarathy, R. A. Rapp, M. Opeka, et al., J. Am. Ceram. Soc. 95, 338 (2012).  31. W. G. Fahrenholtz, J. Am. Ceram. Soc. 90, 143 (2007).  32. V. N. Parfenenkov, R. G. Grebenshchikov, and N. A. Toropov, Dokl. Akad. Nauk SSSR 185, 840 (1969).  33. Thermodynamic Properties  of Materials, Ed.  by V. P. Glushko (Nauka, Moscow, 1978-1982), Vols. 1- 4 [in Russian].  34. A. V. Shevchenko, L. M. Lopato, V. D. Tkachenko, and A. K. Ruban, Inorg. Mater. 23, 225 (1987).  Translated by O. Fedorova  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 59   No. 12   2014  \\x0c']"
},{
  "_id": 80,
  "PDF": "High temperature ablation behavior of pressureless sintered Ta0.8Hf0.2C based ultra-high temperature ceramics.pdf",
  "Text": "['Journal of the European Ceramic Society 40 (2020) 1784-1789  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e l sev ie r .com / loca te / jeu rce ramsoc  Short communication  High temperature ablation behavior of pressureless sintered Ta0.8Hf0.2Cbased ultra-high temperature ceramics  T  Buhao Zhanga,d, Jie Yina,*, Jiaqi Zhenga,d, Xuejian Liua, Zhengren Huanga,b,*, Ján Duszac, Dongliang Jianga  a State Key Laboratory of High-Performance Ceramics and Superfine Microstructure, Shanghai China b Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo, 315201, China c Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, 04353, Košice, Slovak Republic d University of Chinese Academy of Sciences, Beijing 100049, China  Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050,  A R T I C L E  I N F O  A B S T R A C T  Keywords:  Ta0.8Hf0.2C  SiC Double reaction layers Diffusivity of oxygen  Ablation behavior of Ta0.8Hf0.2C and Ta0.8Hf0.2C-10 vol%SiC ceramics was investigated by plasma flame ablation test in air. Linear ablation rate was used for the characterization of ablation resistance and grazing-incidence Xray diffraction (GIXRD), scanning electron microscopy (SEM) and electron backscattered diffraction (EBSD) were used for characterization. Ta0.8Hf0.2C-10 vol%SiC exhibited significantly higher ablation resistance, with ablation rate of 0.7 μm/sec in comparison to the monolithic system with 3.5 μm/sec. An unstable single oxide layer of monolithic Ta0.8Hf0.2C with thickness approximately 20 μm was observed after ablation. Oxidation reactions happened individually as TaC and HfC while ‘self-sacrificial’ volatilization was mainly controlled by active oxidation of TaC. Hf6Ta2O17 was generated by the combination of HfO2 and TaO. A more stable double reaction-layer was observed after ablation of Ta0.8Hf0.2C-10 vol%SiC ceramic. Passive oxidation of TaC was dominant. Textured Ta2O5 skeleton, filled with Hf-containing TaO6, stopped consumption of SiC and hampered oxygen further diffusion.  1.  Introduction  Ultra-high temperature ceramics (UHTCs) have been proposed to serve as thermal protection system (TPS) materials for various applications at temperatures above 2000 °C, such as sharp leading edges and nose cones of next-generation hypersonic vehicles, rocket propulsion lining and advanced energy systems [1,2]. Among the UHTCs, ternary TaxHf1-xC (0 < x < 1) ceramics have attracted worldwide attention in recent decades due to their ultra-high melting point (> 3700 °C) [3] Owing to the same crystal structure with NaCl (B1, space group Fm3m), continuous TaxHf1-xC (0 < x < 1) solid solution could be obtained with TaC and HfC across the whole compositional ranges [4]. Diffusion-driven mass transport mechanism contributes mostly to solid-state densification of TaxHf1-xC (0 < x < 1) ceramics [5]. Since the diffusion activation energy of HfC is higher (HfC: 60.42 kJ/mol, TaC: 39.49 kJ/mol) [6], TaC-rich solid solution carbide (Ta0.8Hf0.2C) is more cost-effective to be consolidated with commercially available TaC  and HfC powders. In our previous work [7], pressurelessly sintering (PLS) of Ta0.8Hf0.2C was favorable for the fabrication of complexshaped components with a relative density of 98.8 %. Besides, Ta0.8Hf0.2C derived from 4TaC-HfC recipe is the most promising candidate as TPS materials for its highest melting point up to 4000 °C among TaxHf1-xC (0 < x < 1) [8]. The ablation resistance ability of UHTCs is not only determined by high melting points, but also dependent on integrity of oxide scale exposed to extreme scouring of heat fluxes [9]. Simulated ablation tests have been conducted on transitional carbide-based UHTCs, including oxy-acetylene flame, plasma flame and plasma wind tunnel. As reported by previous researchers, high bonding strength of Hf-Si-O sticky glassy phase improved the ablation resistance of gradient HfC-SiC ceramic coating of up to 2500 °C with oxyacetylene flame [10]; B-doped Zr-Ti solid solution carbide was also designed and displayed superior ablation resistance during temperature ranges of [2000, 3000]°C under an oxyacetylene flame, owing to the oxide layer of low vaporization rate  ⁎ Corresponding authors at: State Key Laboratory of High-Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China. E-mail addresses: jieyin@mail.sic.ac.cn (J. Yin), zrhuang@mail.sic.ac.cn (Z. Huang).  https://doi.org/10.1016/j.jeurceramsoc.2019.11.043 Received 28 May 2019; Received in revised form 2 October 2019; Accepted 13 November 2019  Available online 16 November 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'B. Zhang, et al.  Journal of the European Ceramic Society 40 (2020) 1784-1789  and oxygen permeability, outstanding self-healing ability and adhesion force with substrate [11]; HfC/TaC-15 vol% MoSi2 ceramics were ablated in a mixture of oxygen/butane/propane at temperatures between 1300 °C-1900 °C with silica pockets found to fill the porosities of oxide scale [12]. Compared to the oxy-acetylene flame and plasma wind tunnel methods, plasma flame ablation technique could provide high enthalpy and plasma flows. A double-layered UHTC coating was designed, containing external ZrC-SiC and inner ZrC-MoSi2 layers, that effectively reduced destructive effect of SiC-depleted zone and bubble burst under plasma flame ablation test [13]; The ablation behaviors of C/SiC-HfC composites were studied at plasma wind tunnel under four different conditions: at low heat flux and stagnation pressure, the HfCSiC composites were oxidized by atomic oxygen to form an reaction layer with a SiO2 top layer and a mixture of HfO2-SiO2 [14]. Possible oxidized product of Ta0.8Hf0.2C, the only ternary crystalline Ta stabilized HfO2 in Ta-Hf-O system: Hf6Ta2O17 [15], would form with non-phase transformation during ablation. A higher content of low melting point Ta2O5 was reported to exacerbate liquid attack on matrix grains during ablation of Ta0.8Hf0.2C sample [16]. The ablation resistance and mechanism of Ta0.8Hf0.2C ceramic is still chemically unclear. Gaseous CO and evaporation of tantalum oxide [17] that can be detrimental to the mechanical integrity of Ta0.8Hf0.2C when exposed to high-velocity and extremely hot heat flux should be taken into consideration as well. SiC with diverse morphologies (eg. equiaxed, platelet-like or whisker) is the most widely used dopant for UHTCs [18], not only for improving the thermo-mechanical properties but for the ablation resistance as well. SiC is expected to improve oxidation-resistance of Ta0.8Hf0.2C effectively [19] by forming a protective oxide layer to inhibit or slow down the oxygen diffusion during ablation [20]. The aim of the present investigation is to study the influence of SiC addition on the ablation resistance of pressureless sintered Ta0.8Hf0.2C based ceramics, to describe the phase composition and microstructure of the reaction layers and to understand the ablation mechanisms in Ta0.8Hf0.2C and Ta0.8Hf0.2C-SiC systems.  2. Experimental procedure  Ta0.8Hf0.2C  Pressurelessly sintered and Ta0.8Hf0.2C-10 vol%SiC ceramics have been used as experimental materials with characteristic microstructure and basic properties, Fig. S1 and Table S1 are in supplementary materials [7]. The ablation resistance of the specimens with a diameter of 30 mm and height of 5 mm, was tested using a plasma flame, generated by a F4-MB gun in air. Surfaces of samples were diamond-polished to a 1 μm finish before exposure to plasma flame for 90 s. The ablation experiment was conducted by an automatic routine: the heat flux was 4.02 MW·m−2, gauged by a Gardon heat flux meter (Shanghai tuxin electronic technology co., Ltd., China). The measured surface temperature was focused on the central area of samples and measured by an optical pyrometer (Marathon MR1SC, Raytek, USA) in the 2-color mode with a spot size of about 2.3 mm (estimated on the basis of the distance from the pyrometer to the sample). The emissivity slope was set as 0.95 for all samples with reference to other reports [21,22]. Linear below:  calculated according  to the  formulas  ablative  rates were  Rl = (dt  d0)/t  Where Rl is linear ablative rate; d0 and dt were the thicknesses of the samples before and after ablation, respectively; t is the ablation time. The phase composition of the reaction layers was identified by grazing incidence X-ray diffraction (GIXRD; Ultima IV diffractometer, Rigaku, Tokyo, Japan) with Cu Ka (k = 1.54056 Å) radiation. The initial surface and cross-section morphologies of ablated samples were investigated by scanning electron microscopy (SEM; Magellan 400, FEI, Hillsboro, American) equipped with Electron Backscattered Diffraction  1785  Table 1  Linear ablation rates after 90 s of ablation, and the highest surface temperature during ablation, of Ta0.8Hf0.2C and Ta0.8Hf0.2C-10 vol% SiC ceramics.  Ta0.8Hf0.2C  Ta0.8Hf0.2C-10 vol%SiC  Thickness before ablation (mm) Thickness after ablation (mm) Linear ablation rate (μm/s) Highest surface temperature (°C)  3.20 2.88 3.5 2098  2.91 2.84 0.7 2579  (EBSD) and energy-dispersive spectroscopy (EDS; Inca, Oxford Instrument, UK). Before cross-section morphology observation the ablated samples were cross sectioned and polished to 1 μm finish.  3. Results and discussion  Ta0.8Hf0.2C  Fig. S1 a,b shows the characteristic microstructure of investigated Ta0.8Hf0.2C and Ta0.8Hf0.2C-10 vol% SiC ceramics before ablation. The average grain size of Ta0.8Hf0.2C system decreased from 13.6 μm to 5.7 μm and relative density increased from 98.8%-99.6% by the addition of 10 vol% SiC. The tested samples retained their original shape after 90 s exposure to plasma flame under air atmosphere. The surface after ablation can be divided into several distinct regions: (I) central (II) transitional and (III) outer ablation regions, [13]. The linear ablation rates after 90 s ablation and at the highest surface temperature of the and Ta0.8Hf0.2C-10 vol% SiC ceramics are listed in Table 1. The high linear ablation rate (3.5 μm/s) in the central area of ablated Ta0.8Hf0.2C with surface temperature up to 2098° indicate a high ablative consumption rate of ceramic matrix. The surface temperature of Ta0.8Hf0.2C-10 vol% SiC during ablation reached 2579 °C, which together with the very low, only 0.7 μm/s linear ablative rate, outline much better ablation-resistance with SiC addition. The central ablated surface and the microstructure of the central cross section of Ta0.8Hf0.2C-based ceramics exhibited different morphologies after ablation, Fig. 1. The morphology of the ablated surface of monolithic Ta0.8Hf0.2C system was composed of looselypacked micro-discs (Fig. 1a), with obvious spalling phenomenon. TaO and Hf6Ta2O17 phases have been detected in the layer according to the GIXRD results (Fig. 1e). TaO (PDF # 19-1299), as a type of tantalum suboxide, has low vapor pressures at high temperature [23]. EDS line the 20 μm thick residual oxide layer rescanning analysis (Fig. 1c) of vealed gradual decrease in Hf and O contents by distance inward, while Ta element displayed a strong inverse tendency. The majority of gaseous TaO, a low vapor pressure phase [24], evaporated during ablation leaving a few condensed TaO on the surface of monolithic Ta0.8Hf0.2C system. Mainly Hf6Ta2O17 phase remained on the surface of ablated oxidation layer based on the GIXRD and EDS results. Unlike those oxidized products of binary carbide (TaC or HfC), the ternary product Hf6Ta2O17, exceeding Ta2O5 and HfO2 for higher melting point and non-phase transformation during elevated temperature stage, was generated from the oxidation of our ternary carbide (Ta0.8Hf0.2C). The ideal mole ratio HfO2/Ta2O5 to form ternary oxidized product without phase transformation should be 6 (6 HfO2 + Ta2O5 = Hf6Ta2O17). However, it cannot be assumed that the formation of Hf6Ta2O17 phase is densely protective upon ablation. The oxidized surface of Ta0.8Hf0.2C-10 vol%SiC ceramic was more porous after ablation with higher average pore size from vertical view in comparison to the reaction layer of the monolithic system but without of distinct spalling phenomenon, (Fig. 1b). Cross-sectional view - of thicker ablated (oxidized) layer revealed a presence of double layer, which consisted of outer (60 μm) and inner (80 μm) layers, both formed by Ta and Hf oxides. Damage of outer oxidation layer (bubble burst) was observed and the inner layer, generated from the inward diffusion of oxygen, contained internal unreacted SiC particles,  \\x0c', 'B. Zhang, et al.  Journal of the European Ceramic Society 40 (2020) 1784-1789  Fig. 1. Surface and cross-sectional morphologies of ablated samples: (a)&(c) Ta0.8Hf0.2C and (b)&(d) Ta0.8Hf0.2C-10 vol%SiC ceramics, line scanning of cross-sectional samples; GIXRD patterns of ablated (e) Ta0.8Hf0.2C and (f) Ta0.8Hf0.2C-10 vol%SiC ceramics.  inserted in (c)&(d) are EDS  Fig. 1d. Si element signal appeared only at specific equiaxed SiC grains of inner layer from EDS-line scanning inserted in Fig. 1d. Owing to the high vapor pressure of SiO (g) under reducing conditions (R4), active oxidation of SiC under the oxide scale produced a SiC-depleted region [25]. Triclinic Ta2O5 (PDF#21-1198) and Hf6Ta2O17 (PDF#44-0998) phases were detected as the ablated products based on GIXRD patterns of Fig. 1f. No obvious Si-containing solid-state phase was shown in GIXRD patterns. Residual SiC particles embedded inside the inner layer contribute any signals due to the limited penetration depth of GIXRD at grazing angles [26]. Cross-sectional microstructure observation (together with EBSD analysis) of the ablated layers of Ta0.8Hf0.2C-10 vol%SiC ceramic revealed phases of Ta2O5 (yellow color, outer and inner layers), TaO6 (red and α-SiC (blue color, outer and inner layers) color, inner layer), Fig. 2a. Ta2O5 grains in the outer ablated layer exhibited a textured and rod-like morphology, while α-SiC particles survived only within the inner layer. Combined with GIXRD result above, it can be confirmed that the active oxidation of SiC happened and lead to the SiC-depleted  region of outer ablated layer. With both outer and inner layers inhibiting the diffusion of oxygen, impressive decrease in oxygen content was hence found from inner layer to the matrix (Fig. 2d). As shown in Fig. 3a, the structure of Ta2O5 (in yellow) could be high temperature phase (H-Ta2O5) with lattice parameters of a = 3.86A˚ and c = 36.18A˚ that was reported to own a liquidus temperature of 1877 ± 40 °C [27]. The EDS maps, Fig. 3 (b), were conducted inside the red rectangle area. As known, the EBSD signal is very sensitive to surface roughness. Unlike the EDS result within the larger area, EBSD result only within the narrow strip area (black rectangle area, which was smooth) was more precise, and inserted in Fig. 3 (a). EBSD analyses revealed that in outer ablated layer Ta2O5 grains are present mainly with (100)-orientation, Fig. 3c. It indicated that the Ta2O5 grains underwent a process of crystallization with preferential growth toward along the c-axis direction, driven by gradient temperature from outermost ablation area (up to 2579 °C). Quite recently, Scott J et al. [28] reported the disordered crystal structure of Hf6Ta2O17 with one set of 6oxygen-coordinated cation formed an octahedron. It was implied from  1786  \\x0c', 'B. Zhang, et al.  Journal of the European Ceramic Society 40 (2020) 1784-1789  Fig. 2. Detailed cross-sectional morphologies of the ablated Ta0.8Hf0.2C-10 vol%SiC specimen: (a) overall view of the outer layer, inner layer, and matrix, with EBSD phase-map inserted, (b) outer ablated layer and (c) inner ablated layer; (d) O-element EDS line scanning of red line in (c) from inner layer to the ceramic matrix.  Fig. 3. (a) EBSD phase-map, superimposed on the outer ablated layer; (b) EDS mapping of inserted red rectangle area in (a), yellow, green and red colours correspond to tantalum, hafnium and oxygen; (c) pole figures of textured Ta2O5 in the ablated Ta0.8Hf0.2C-10 vol%SiC surface.  1787  \\x0c', 'B. Zhang, et al.  Journal of the European Ceramic Society 40 (2020) 1784-1789  oxygen further diffusion. The active oxidation of TaC (R1) was suppressed while the oxidation of SiC (R5) occurred. As shown in Fig. 4, Gibbs free energy of active oxidation (R1) and passive oxidation (R2) reactions of TaC revealed the preference of oxidation reactions differed by temperatures: higher temperature was required for active oxidation (R1) of TaC. The incorporation of SiC (thermal conductivity: 180 W·(m·K)−1) enhanced the heat conduction by dissipating the heat and lowering the oxidation kinetics [29]. The surface temperature could be heavily influenced by the effective heat resistance of the specimen, e.g. thermal diffusivity and thickness. However, the difference between the thicknesses of samples with and without SiC addition in our research was not obvious, as shown below in Table 1. Besides, the thermal diffusivity and conductivity of monolithic ceramic are 7.2 mm2∙s−1 and 18.6 W·(m∙K)−1 respectively at room temperature, lower than sample with SiC addition (11.4 mm2∙s−1 and 26.7 W·(m∙K)−1) according to Table S1 from our supplementary materials. The higher ablation rate of monolithic Ta0.8Hf0.2C ceramic indicated that its bulk matrix was ablated rapidly. Active oxidation of TaC (R1; it should be noted that R1 represented the active oxidation of TaC, while R2 represented the passive oxidation of TaC) dominated and gaseous TaO was formed for monolithic Ta0.8Hf0.2C ceramic. Passive oxidation (R2) product (Ta2O5) was obtained after the ablation of Ta0.8Hf0.2C -SiC composite. The vaporization of TaO from R1 is probably responsible for the lowered surface temperature (2098 °C) of the monolithic Ta0.8Hf0.2C sample during ablation, when comparing to that (2579 °C) of Ta0.8Hf0.2C-10 vol%SiC composites. Homogenous oxidation of remaining Ta0.8Hf0.2C occurred inside gaps between textured Ta2O5 particles. Hf/Ta-containing oxides connected Ta2O5 particles, resisting high speed airflow and impeding the oxygen diffusion to inner non-oxidized region as well.  Ta0.8Hf0.2C  4. Conclusions  The aim of present investigation was to study influence of SiC addition on ablation behavior of pressureless sintered Ta0.8Hf0.2C-based ceramics. The main results are following:  • Ta0.8Hf0.2C-SiC ceramic exhibits significantly higher ablation resistance with ablation rate of 0.7 μm/sec in comparison to the monolithic system with ablation rate of 3.5 μm/sec. • the residual single oxide layer of ablated monolithic system with approximately 20μm was thickness composed of condensed tantalum suboxide TaO and Hf6Ta2O17 phases. An oxidation mechanism of Ta0.8Hf0.2C was proposed with dominate active oxidation of TaC, together with the oxidation of HfC individually. The high ablation rate and the low thickness of the layer are the evidence that the layer is moving fast inward and consuming the bulk • A double ablation-layer was observed after ablation of Ta0.8Hf0.2Cmatrix rapidly. SiC ceramic, with a porous and SiC-depleted outer layer, supported by textured Ta2O5 gains and with a dense multi-phase inner layer containing survived SiC particles. Thanks to the addition of SiC with high thermal conductivity the ablation resistance of Ta0.8Hf0.2C was enhanced by dissipating heat and lowering the oxidation kinetics. The active oxidation of TaC was suppressed and oxidation of SiC occurred instead, which benefited anti-ablation behavior. This double-layer insured extremely high ablation protection at this very high testing temperature.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Fig. 4. Thermodynamical analysis of possible reactions during the ablation test of theTa0.8Hf0.2C-SiC ceramics.  EDS and EBSD mappings that Hf-containing TaO6 octahedra oxides with fine grain size (beyond EBSD detection limit) are filling the gaps around the textured rod-like Ta2O5 in the outer layer. During the ablation of Ta0.8Hf0.2C-SiC system the following oxidation reactions can occur. As shown in Fig. 4, all reactions would happen at above 1000 °C.  TaC+O2(g) = TaO(g) + CO(g)  2TaC + 7/2O2(g) = Ta2O5(l) + 2CO(g)  HfC + 3/2O2(g) = HfO2(s) + CO(g)  SiC + 3/2O2(g) = SiO2(l) + CO(g)  SiC+O2(g) = SiO(g) + CO(g)  (R1)  (R2)  (R3)  (R4)  (R5)  Ta0.8Hf0.2C  Timely release of CO was a distinct feature for Ta0.8Hf0.2C-based ceramics. Two possible oxidation pathways for TaC, as comparable to the active oxidation and passive oxidation routes of SiC, corresponded to oxidation products during these experiments: gaseous TaO generated via active oxidation (R1) of monolithic Ta0.8Hf0.2C while passive oxidation (R2) product (liquid Ta2O5) was obtained after the ablation of Ta0.8Hf0.2C-SiC. Under the present ablation condition with extreme temperature (higher than 2000 °C), SiO2 would be formed in liquid phase (R4), or even in gaseous state (R5). Based on the ablation results and microstructure characteristics of ablated layers detailed ablation mechanisms of based ceramics can be proposed as follows: For ablation mechanism of monothetic Ta0.8Hf0.2C solid solution, oxidations happened individually as TaC and HfC. Our results show that active oxidation (R1) of TaC occurred inside Ta0.8Hf0.2C solid solution predominantly. Gaseous TaO vanished after ablation, leaving a few condensed sub-stoichiometric TaO and forming Ta-Hf-O oxide at the surface of monolithic Ta0.8Hf0.2C. As a thorough solid solution carbide Ta0.8Hf0.2C, the diminishment of oxidation layer conducted homogenously without any TaC-depleted region. The ‘self-sacrificial’ active oxidation behavior of TaC is detrimental, together with the oxidation HfC, that desintergrate the solid solution phase in O-containing atmosphere. Hf6Ta2O17 was further generated by the combination of HfO2 and deposited (Ta,O) upon the oxidation of HfC. For ablation mechanism of Ta0.8Hf0.2C-10 vol%SiC, the active oxidation of TaC was suppressed with SiC addition. As reported in Liu et al.’s work [13], SiC depletion occurred for ZrC-SiC ceramic coating similarly as it was observed in our experiment. Passive oxidation mode (R2) of TaC determined a textured Ta2O5 skeleton filled with possible Hf6Ta2O17 on the porous outer ablated layer. The consumption of SiC stopped in the inner ablated layer as Ta2O5 and Hf6Ta2O17 particles connected with each other to resist high speed airflow and hamper  1788  \\x0c', 'Journal of the European Ceramic Society 40 (2020) 1784-1789  [13]  [16]  [20]  [11]  [12]  [17]  [18]  [14]  [15]  tantalum at 1300-1800°C, J. 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Han, Z. Zhang, L. Ji, Y. Jiang, The crystal structure of high temperature phase Ta2O5, Acta Mater. 55 (7) (2007) 2385-2396. S.J. McCormack, R.J. Weber, W.M. Kriven, In-situ investigation of Hf6Ta2O17 anisotropic thermal expansion and topotactic, peritectic transformation, Acta Mater. 161 (2018) 127-137. [29] G.A. Slack, Thermal conductivity of pure and impure silicon, silicon carbide, and diamond, J. Appl. Phys. 35 (12) (1964) 3460-3466.  [21]  [22]  [23]  [26]  [28]  B. Zhang, et al.  Acknowledgements  Financial support from National Natural Science Foundation of China (No. 51602325), Youth Innovation Promotion Association (CAS, No. 2018289), Science Foundation for Youth Scholar of State Key Laboratory of High-Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics CAS (SKL201602), Scientific and Technological Innovation Project of Shanghai Institute of Ceramics are gratefully acknowledged. Buhao Zhang would also like to acknowledge the Chinese Scholarship Council (CSC) for financial support.  Appendix A. Supplementary data  Supplementary material related to this article can be found, in the online version, at doi:https://doi.org/10.1016/j.jeurceramsoc.2019.11. 043.  References  [1] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for extreme environments, Scripta Mater. 129 (2017) 94-99. [2] X. Zhang, P. Hu, J. Han, S. Meng, Ablation behavior of ZrB2-SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions, Compos. Sci. Technol. 68 (7-8) (2008) 1718-1726. [3] O. Cedillos-Barraza, D. Manara, K. Boboridis, T. Watkins, S. Grasso, D.D. Jayaseelan, R.J. Konings, M.J. Reece, W.E. Lee, Investigating the highest melting temperature materials: a laser melting study of the TaC-HfC system, Sci. Rep. 6 (2016) 37962. [4] O. Cedillos-Barraza, S. Grasso, N.A. Nasiri, D.D. Jayaseelan, M.J. Reece, W.E. 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},{
  "_id": 81,
  "PDF": "High temperature creep of 20 vol_. SiC-HfB2 UHTCs up to 2000 °C and the effect of La2O3 addition.pdf",
  "Text": "['Journal of the European Ceramic Society 38 (2018) 47-56  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e l sev ie r .com / loca te / jeu rce ramsoc  Feature article  High temperature creep of 20 vol%. SiC-HfB2 UHTCs up to 2000 °C and the effect of La2O3 addition  MARK  E. Zapata-Solvasa,b,⁎, D. Gómez-Garcíaa,c, A. Domínguez-Rodríguezc, W.E. Leeb  a Instituto de Ciencia de Materiales de Sevilla, CSIC-Universidad de Sevilla, 41092 Sevilla, Spain b Centre for Nuclear Engineering & Dpt. of Materials, Imperial College London, SW7 2AZ, UK c Dpto. Física de la Materia Condensada, Universidad de Sevilla, 41012 Sevilla, Spain  A R T I C L E  I N F O  A B S T R A C T  High temperature compressive creep of SiC-HfB2 UHTCs up to 2000 °C has been studied. Microstructural analysis after deformation reveals formation of new phases in the Hf-B-Si and Hf-B-Si-C systems, which are responsible for the poor creep resistance. RE oxide additions have a negative effect reducing the creep resistance of SiC-HfB2 UHTCs. A simplistic analysis for the required creep resistance is described, indicating that only SiC-HfB2 UHTCs could withstand re-entry conditions for 5 min in a single use. However, RE oxide addition to SiC-HfB2 UHTCs does not provide the required creep resistance for them to be candidate materials for hypersonic applications.  Keywords: Ultra-high temperature ceramics Creep Cavitation damage Reaction damage Limited ductility  1.  Introduction  Me borides and carbides (Me = Zr or Hf) with melting points in excess of 3000 °C are promising candidates for thermal protection systems (TPS) in hypersonic vehicles, including sharp-leading edges (SLE) and sharp-nose cones (SNC) [1,2], and belong to a family of materials termed ultra-high temperature ceramics (UHTCs). Ideally, UHTCs components for hypersonic applications have to maintain a stable structure under the operating conditions, which involve no phase transformations and no shape changes during operating conditions. No shape changes require high creep resistance and, e.g the components should deform no further than the elastic limit to be reusable. In addition, UHTCs components require a combination of properties at high temperatures, such as high temperature strength, oxidation resistance, with high thermal conductivity being particularly desirable to maximize thermal transport during exposure in high-temperature reactive environments [3]. Me borides possess a better combination of oxidation resistance and thermal conductivity than Me carbides and have been the subject of intense research over the last 15 years [4,5]. Moreover, ZrB2 and HfB2 exhibit higher thermal conductivity (100 W/(m K)) at room temperature (and > 50 W/(m K) at 1900 °C) [6] than their respective carbides (< 40 W/(m K) from room temperature to 800 °C) [5]. However, the oxidation resistance of MeB2 is relatively poor which could compromise the structural stability of UHTCs components under hypersonic operating conditions. As a consequence, UHTCs research has focused over the last decade on the oxidation resistance of MeB2based UHTCs and different approaches to improve oxidation resistance  have been attempted; (i) Adding different Si-containing compounds to incorporate different elements in the outer protective borosilicate (BS) coating, which results in an increase of viscosity and melting temperature as well as a reduction in the oxygen diffusion coefficient through the BS [7]; (ii) Adding different diborides, such as TaB2, TiB2 or CrB2, as BS glasses containing oxides of the elements listed are immiscible leading to phase separation, increasing their viscosity and melting temperature [8]; (iii) Adding rare earth compounds that could lead to a protective refractory coating formed during the oxidation of the UHTCs components [9,10]; (iv) Adding low contents of rare earth compounds (< 3 vol%), such as La2O3, which lead to the formation of MeOxCy particles in the oxide scale as a result of increasing viscosity of BS, leading to a long-term stabilization of the oxide scale [11]. The composition that has received most attention to date and is currently considered as a baseline UHTCs composition for hypersonic TPS development is SiC-reinforced UHTCs, with a SiC content ranging from 10 to 30 vol.% [2,9]. Oxidation of SiC produces an outer BS protective glass that could protect UHTCs components up to 1650 °C, at which temperature pure SiO2 melts and bulk UHTCs would be exposed to continuous oxidation. Although the BS coating melts at lower temperatures it could withstand temperatures up to 2000 °C for up to [12-17]. However, 1 h under static oxidation conditions this liquid layer will be blown away during a hypersonic reentry. On the other hand, hypersonic applications require high strength at high temperature and estimated conditions for hypersonic re-entry are 1800 °C and 400 MPa [3]. Furthermore, the latter conditions are at the less harsh end of reentry conditions and in the long term, the  ⁎ Corresponding author at: Centre for Nuclear Engineering & Dpt. Of Materials, E-mail address: eugenio.zapata-solvas@imperial.ac.uk (E. Zapata-Solvas).  Imperial College London, SW7 2AZ, UK.  http://dx.doi.org/10.1016/j.jeurceramsoc.2017.08.028 Received 27 April 2017; Received in revised form 16 August 2017; Accepted 19 August 2017  Available online 24 August 2017 0955-2219/ © 2017 The Author(s). Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/BY/4.0/).  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  development of UHTCs able to withstand temperatures in excess of 2000 °C under the same stresses will be required. This requirement has led the scientific community to develop high strength UHTCs components, as high as 1 GPa at room temperature (RT) [18,19]. Moreover, SiC reinforced MeB2 have shown the highest RT strength values [20]. However, UHTCs must withstand extreme conditions of temperature and stress under harsh environments and the available data about high temperature mechanical properties, i.e. strength and creep resistance, is scarce. Concerning HT strength measurements, there have been few reports of the strength of UHTCs at the expected operating temperatures. ZrB2 shows a strength value of 220 MPa in the temperature range from 1600 °C to 2200 °C [21]. Hu and Wang reported a strength of 220 MPa for ZrB2-30 vol.% SiC at 1800 °C [22] while Neuman et al. found a strength of 540 MPa at 1800 °C and 260 MPa at 2200 °C in ZrB2-30 vol.% SiC + 2 wt.% B4C illustrating the impact on strength of a small amount of B4C [23]. Highest strength at high temperature till date, 290 MPa at 2200 °C and 2300 °C, was reported by Neuman et al. in ZrB2-10 vol.% ZrC [24]. However, the strength of ZrB2-10 vol.% ZrC at 1800 °C was 345 MPa which is lower than the 540 MPa of ZrB2-30 vol.% SiC + 2 wt.% B4C at 1800 °C. All the strengths mentioned above were measured in inert atmosphere. However, measured strengths in air for ZrB2-30 vol.% SiC + 2 wt.% B4C were 360 MPa at 1500 and 1600 °C [25], much lower than the 540 MPa shown at 1800 °C in Ar [23]. The main limitation for high temperature mechanical properties testing at 2000 °C is that the majority of commercial furnaces are limited to 1600 °C, especially those furnaces able to test in air, and there are few facilities around the world with the capability of testing at 2000 °C under inert atmospheres. Those strengths are below the expected stress conditions mentioned above, which have been predicted considering elastic deformation of UHTCs components. Therefore, the study of active deformation mechanisms, which could produce a stress relaxation during hypersonic re-entry to avoid mechanical failure of UHTCs, becomes crucial. Furthermore, the structure of hypersonic vehicles depends on the structural integrity of the UHTCs components, which is directly related to their creep resistance. High creep resistance is desirable to avoid shape changes during plastic deformation. For example, the edge of SLEs and SNCs are the areas of UHTCs components exposed to highest temperature and maximum stresses, which are susceptible to plastic deformation under the operating conditions. Therefore, a small change in the radius of curvature of SLE or SNC during hypersonic re-entry could have a significant negative impact on maneuverability and subsequent vehicle aerodynamics. Few creep studies have been carried out on UHTCs, with most that have being for SiC-reinforced UHTCs. All have been performed in flexure [22-26] and for deformations < 1% [26-29]. Deformation in flexure exceeding 2% leading to mechanical failure in 20 vol.% SiCreinforced ZrB2 [28], has been examined by Guo et al., who observed formation of isolated cavities in the tensile side of 30 vol.% SiC-reinforced ZrB2, nucleated always on SiC particle boundaries [28] for deformations as low as 0.2% at 1500 °C under strain rates of 2 × 10−9 s−1. Talmy et al. observed crack formation during the tertiary stage of creep in the tensile side of 50 vol.% SiC-reinforced ZrB2, presumably related to cavity nucleation, and determined that the controlling deformation mechanism is SiC grain boundary sliding [27]. Bird et al. reported formation of cavities in 20 vol.% SiC-reinforced ZrB2 above 1600 °C and suggested that cavitation partially accommodates grain boundary sliding [29]. In addition, it was determined that diffusional creep accommodated by lattice diffusion is the active deformation mechanism below 1600 °C and grain boundary sliding above 1600 °C [29]. Note that true creep behavior is related to flexural creep in the case of symmetric creep, i.e. in which the tensile creep rate is the same as that in compression. Gangireddy et al. suggested that 30 vol.% SiC-reinforced ZrB2 has symmetric creep for strains below 1% [26] based on transmission electron microscopy (TEM) observation of the absence of glassy films at grain boundaries in 10 vol.% SiC-reinforced ZrB2 fabricated by the same hot press (HP) method, carried out by  48  Jayaseelan et al. [30]. However, Bird et al. quantified a neutral axis shift during deformation in 20 vol.% SiC-reinforced ZrB2 which confirms the existence of asymmetric creep [31]. The asymmetry between tensile and compressive creep is a typical feature of SiC-, Al2O3or Si3N4-based ceramics, in which an easily deformable glass is present at the grain boundaries showing an accelerated creep under tension [32,33]. In addition, the presence of cavitation during plastic deformation as observed in reaction bonded SiC is evidence of asymmetric creep behavior [33]. Cavitation usually leads to asymmetric creep behavior as coalescence of cavities is easier under tensile than compressive creep [34]. Monolithic ceramics, such as Y-TZP and Al2O3, show symmetric creep behavior [32,35]. Cavitation in tensile or flexure creep makes the assessment of active deformation mechanisms more difficult as deformation mechanisms are usually the combination of at least two deformation mechanisms [33,34]. Therefore, deformation mechanisms are usually studied in compression by the ceramics community [36,37] unless the stresses are low enough to have pure diffusional creep without cavitation [34]. Moreover, the symmetric creep behavior of SiC-reinforced ZrB2 ceramics is questionable as there are no data under tensile or compressive creep and the presence of cavities after deformation has been detected only on the tensile side. The work described in this paper is the first to examine compressive creep of 20 vol.% SiC-HfB2 ceramics in an attempt to identify the active deformation mechanisms from 1800 °C-2000 °C, which is the expected temperature range for SiC-reinforced MeB2 UHTCs under hypersonic regimes. Compressive creep is analyzed in terms of a constitutive creep equation [37];  ε˙  =  A Gb kT  b d  ⎛ ⎝  ⎞ ⎠  p  σ G  ⎛ ⎝  n  ⎞ ⎠  D exp  0  Q kT  ⎞ ⎠  −  ⎛ ⎝  (1)  where ε˙ is the strain rate, A dimensionless constant, G the shear modulus, b the magnitude of the Burgers vector or any characteristic length scale in the material, k the Boltzmanńs constant, T the absolute temperature, d the average grain size and σ the applied stress. The term D0 is the frequency factor of an appropriate diffusion coefficient responsible for the migrating species involved in the accommodation process. Q is the activation energy of the active diffusion phenomena. A deformation mechanism is usually identified by the determination of p, n, Q and through the observation of microstructure after deformation. Recently, it was reported that the addition of 2 wt.% La2O3 to SiCreinforced MeB2 was beneficial for long-term oxidation processes as the oxide layer showed better stability than SiC-reinforced MeB2 due to MeOxCy formation within the oxide layer [38]. Therefore, the aim of this study is to characterize the structural properties and plastic response of 20 vol.% SiC-reinforced HfB2 and 2 wt.% La2O3 + 20 vol.% SiC-reinforced HfB2 under compression from 1800 °C to 2000 °C and determine whether SiC-reinforced HfB2 UHTCs are structurally stable under such conditions.  2. Experimental procedure  d50  5.0 μm, ρ = 10.5 g/cm3, HfB2 powder (> 99%, Sigma (α-SiC, d50  0.7 μm, Aldrich, Gillingham, UK), SiC powder 99%, ρ = 3.217 g/cm3, Good Fellow Chemicals, Huntingdon, UK) and La2O3 d50  10 μm, ρ = 6.51 g/cm3, (> 99%, Fluka chemicals supplied through Sigma Aldrich, Steinheim, Germany) were used to form different UHTC compositions; HfB2 + 20 vol.% SiC (HS20) and HfB2 + 20 vol.% SiC + 2 wt.% La2O3 (HS20La). Ceramic powders were processed and then sintered using Spark Plasma Sintering (SPS) as described previously [10]. The main microstructural features of SPS UHTCs [11,19] are; (I) high density (> 99% of theoretical relative density); (II) HfB2-based UHTCs contain an inhomogeneous dispersion of SiC the HfB2 matrix with SiC agglomerates as large as 20 μm; throughout and (III) La2O3 particles are often in close proximity to SiC particles and improve the SiC dispersion throughout the HfB2 matrix. SPS (40 mm dia × 5 mm thick) billets were cut  electro using  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  discharge machining (EDM) to obtain 2.5 mm × 2.5 mm × 4.5 mm bars, all bar surfaces were ground using a 1200 grit diamond platen to remove surface damage to leave 2 mm × 2 mm × 4 mm final dimensions. Plastic deformation under compression at 1500 °C was studied in a mechanical frame (Zwick/Roll, Munich, Germany) equipped with a furnace with MoSi2 heating elements. Plastic deformation under compression from 1800 °C to 2000 °C was studied in a mechanical frame (Instron, Norwood, USA) equipped with a furnace with tungsten heating elements (GT Technologies, New Hampshire, USA) operated at a pressure of approximately 9 × 10−3 Torr. Cross head speeds ranged from 5 μm/min to 100 μm/min. Push rods were protected by the use of SiC spacers, which are more creep resistant than materials tested in this work. Cross sections of specimens after compressive testing for scanning electron microscope (SEM) characterization were prepared using conventional methods involving successive steps of grinding and polishing with diamond slurries and cloths embedded with up to 1 μm diameter particles. The specimens were observed in an SEM equipped with a field emission gun (Hitachi S5200, Hitachi, Tokyo, Japan). Samples were examined in secondary electron (SE) imaging mode and backscattered electron (BS) image mode for atomic number contrast. Energy dispersive X-ray spectroscopy (EDS) was used to aid phase identification. Specimens for transmission electron microscopy (TEM), equipped with a LaB6 gun (CM200, Philips, Eindhoven, Holland), were prepared using conventional methods of cutting, dimpling and ion milling. The TEM was equipped with an EDAX EDS detector (model R-TEM CM200, Mahwah, New Jersey, USA). Bright-field (BF) imaging and EDS were used to characterize the specimen microstructures and identify the different phases. Genesis spectrum TEM quantum materials software from EDAX was used to carry out the quantification of different elements during phase identification studies. Elemental C content analysis on powders found in HS20La after compressive deformation at 1800 °C were analysed by an elemental analyser (LECO CHNS 932, St. Joseph, Michigan, US). The strain rate range studied was estimated, at a first approximation, as the thermal strain rate related to a heating rate in the range from 200 °C/min to 1000 °C/min, which could be a realistic heating rate range for hypersonic conditions, relating the thermal expansion to the thermal strain according to Eq. (2);  ε˙  =  α E  Δ T Δ t  (2)  where α is the coefficient of thermal expansion (CTE), E the elastic modulus and ΔT/Δt the heating rate. Heating rates of 200 °C/min and 1000 °C/min correspond to thermal strain rates of 2.8 × 10−5 s−1 and 1.4 × 10−4 s−1, respectively, at an elastic modulus of 500 GPa [4] and CTE previously reported [6], which are in the range of interest for hypersonic applications.  3. Results and discussion  3.1. HS20 high temperature mechanical properties  Fig. 1 shows the creep response of HS20 from 1500 °C to 2000 °C approximately over the strain rate range 2.8 × 10−5 s−1 and 1.4 × 10−4 s−1 for strains as high as 24%. At 1500 °C (Fig. 1a), HS20 shows a limited ductility (< 6%) for strain rates ε˙ < 6 × 10−5 s−1. At higher strain rates, the deformation is increased to nearly 12% and significant softening is also observed. Deformation-induced softening could be related to either cavitation or phase instabilities. In addition, true steady states of deformation are only observed at 1900 °C or higher temperatures as softening is still active at 1800 °C. Stresses for the steady states of deformation were calculated through extrapolation of steady state to the initiation of the plastic regime at 1900 °C and 2000 °C, as plotted in Fig. 2a. Calculated stress exponents for the steady states of plastic deformation are shown in Fig. 2a, ranging from 3 to 1 as the temperature increases from 1500 °C to 2000 °C. In addition,  calculated activation energies are 242 and 334 kJ/mol in the temperature ranges 1500-1800 °C and 1900-2000 °C respectively, suggesting the activation of different deformation mechanisms as the stress exponent also changes from 3 to 1 in these same temperature ranges. The last deformation state in Fig. 1c and 1d is not a true stationary as some softening is observed. Therefore, the stress exponent for this last deformation state might not be accurate and this state is not considered for the plot of stationary strain rates versus stress, so the stress exponent shown in Fig. 2a is more accurate than the one displayed for the last deformation state in Fig. 1c and 1d respectively. In addition, no true stationaries were observed at 1800C either, but the data points were considered to allow comparison with previously reported data. The satisfactory agreement of the calculated activation energy value with literature data in Fig. 2b suggests that it was a reasonable assumption. Fig. 3 compares some of the previously reported data for the stationaries of plastic deformation on UHTCs. Studies at a strain rate below 10−5 s−1 [27-29] are not considered as they are out of the range of interest of hypersonic applications. HS20 is between 2 and 3 times more creep resistant than the ZS30 (ZrB2 + 30 vol.% SiC) studied by Gangireddy et al. [22]. The different response could be attributed to the more refractory nature of HfB2 compared to ZrB2 or the lower SiC volume content of HS20 (20 vol.%) compared to ZS30 (30 vol.%). Talmy et al. studied the plastic response of ZrB2 with SiC contents ranging from 0 to 50 vol.% and observed a change in the activation energy from 130 to 509 kJ/mol, showing a plastic behavior characterized by a lower SiC content giving a higher creep resistance. However, highly pure SiC is known to be extremely creep resistant and hard to sinter also, if no impurities or additives are present. Liquid phase sintering (usually from melting of silica contamination) or boron and carbon doping are common approaches in SiC to promote full densification [39-42] and temperatures of 2200 °C required to sinter pure avoid the use of polycrystalline SiC by SPS [43]. For example, Fig. 2b reveals that additive-free nano β-SiC is 2-3 orders of magnitude more creep resistant at 1800 °C than HS20 at 1500 °C. In addition, liquid phase sintered SiC at 1756 °C is 1-2 orders of magnitude more creep resistant than HS20 at 1500 °C. Therefore, SiC is not playing any reinforcing role, or providing any improvement on the creep resistance, from a structural point of view which agrees with observations from Talmy et al. [27], who determined that the higher the SiC content, the higher the activation energy and the softer/lower the creep resistance. Fig. 2b shows activation energy versus SiC content calculated by Talmy et al. for ZrB2 in which the activation data for this study are included. Note that Talmy et al.’s study is from 1200 to 1500 °C. Therefore, the activation energy used to make the comparison is the one obtained in this study in the lower temperature range studied (1500-1800 °C). Good agreement between our estimated activation energy and the activation energies of Talmy et al. [27] is shown in Fig. 2b, suggesting that SiC content might have a stronger influence on the plastic response of MeB2 (Me = Zr, Hf) UHTCs than the kind of MeB2 analyzed. This statement is supported by the plastic response of 1 wt.% boron-doped SiC in Fig. 2b at 1758 °C, which is similar to the plastic response of HS20 at 1800 °C, more creep resistant than ZS30 at 1800 °C, and shows a strong influence of B in the high temperature behavior of SiC. Liquid-phase sintered SiC and additive-free nano β-SiC are more than 3 orders of magnitude more creep resistant than 1 wt.% B-doped SiC at the same temperature (1750 °C) [39,40,43], highlighting the enormous enhancement of creep rate or diffusion rate during plastic deformation by an addition of 1 wt.% B. There is a lack of data about plasticity of monolithic MeB2 to confirm the role of SiC, but Melendez-Martinez et al. studied the plastic response under compressive creep of ZrB2 from 1400 to 1600 °C [44]. In addition, it was reported that at 1500 °C ZrB2 deforms at 10−7 s−1 under 400 MPa, which is 300 times less resistant than HS20 reported herein in spite of ZrB2 having a porosity of 13.5 vol.% [44]. Considering that there is little difference between the plastic response of ZS30 and HS20, as observed in Fig. 3, and also that the activation energy for ZrB2 increases with the addition of SiC, it seems likely that the addition of  49  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  Fig. 1. Stress versus strain for HS20 at; a) 1500 °C, b) 1800 °C, c) 1900 °C and d) 2000 °C. Strain rates (ε˙ ) for each deformation state and subsequent stress  exponents (n) are shown.  SiC degrades the creep resistance of MeB2 UHTCs. Furthermore, it indicates the need to characterize and understand the plasticity of MeB2. Concerning the high temperature creep properties of MeB2 reinforced with MeC, there are a couple of studies in flexure for ZrB2-20 vol.% ZrC and TiB2-20 vol.% TiC in addition to other compositions with MeC contents in excess of 40 vol.%, which are far from the interest of this study [45,46]. ZrB2 + 20 vol.% ZrC deforms at 2000 °C in the stress range 5-30 MPa with a strain rate range of 1.3-8.7 × 10−3 s−1 respectively and a stress exponent of n=1, suggesting grain boundary sliding as the deformation mechanism. However, the ZrB2-20 vol.% ZrC is nearly 100-150 times less creep resistant than the ZrB2-20 vol.% ZrC studied by Gangireddy et al. [26] and HS20 from this study. In addition, contamination during milling of up to 6 wt.% with WC-Co was reported, which could explain the poor creep response of ZrB2-20vol.% ZrC studied by Kats et al. [45] as Co melts at 1500 °C. This example illustrates how important it is to minimize contamination as undesired additional second phases can have a detrimental effect on physical properties at temperatures of 2000 °C or higher. Furthermore, monolithic ZrB2 and TiB2 were also studied at temperatures higher than 2000 °C [45,46]. However, if an extrapolation is made to 2000 °C, they are 100 times less creep resistant than ZS30 in the best case, which highlights the importance of controlling second phase impurities as well as the need to obtain reliable data on plasticity of monolithic MeB2.  Fig. 3. Comparison of  the strain rate versus stress from HS20 with previously reported  data.  3.2. HS20 microstructural characterization after deformation  The concept of SiC-reinforced MeB2 comes from the exceptional strength shown by SiC-ZrB2 composites, as high as 1 GPa [18], and the increased fracture toughness of MeB2 with SiC additions, from  Fig. 2. a) Strain rate versus  stress  for  the deforma tion stationary states of HS20. b) Comparison of the  activation energy obtained in this study for the temperature range 1500-1800 °C with Talmy et. al  study [27].  50  \\x0c', 'E. Zapata-Solvas et al.  Fig. 4. XRD of the external surface of HS20 after deformation at 1800, 1900 and 2000 °C.  XRD before deformation looks identical  than after deformation at 1800 °C.  values of 3 for monolithic MeB2 to values as high as 7 MPa/m1/2 for composites with SiC [19]. However, this trend reverses at high temperature as SiC acts as a softening agent. A possible reason for this softening is the formation of an eutectic in the Me-B-C ternary system with a melting point of 2390 °C for Zr-B-C, which is nearly 1000 °C lower than the melting temperature of ZrB2 [47]. In addition, the Me-BC phase diagrams reported by Rudy and Windisch [48] suggestthat the combination of MeC + MeB2 might be of interest at temperatures lower than 1500 °C as there could be temperature-dependent B-C exchange reactions associated with volume changes in the carbide particles, resulting in bodies with poor mechanical properties and subsequently limiting their high temperature potential use at temperatures above 1500 °C. If we consider the results observed here and previous available data [26,39,40,43], including the effect of 1 wt.% B in SiC, it could be concluded that B also has a negative effect on the high temperature mechanical behavior of MeB2 + SiC systems. Fig. 4, shows XRD patterns of HS20 external surfaces before and after deformation revealing that at temperatures higher than 1800 °C, the formation of HfB is detected at the surface. However, XRD from an internal surface in the bulk does not show formation of HfB at any temperature. Plastic deformation is a phenomenon controlled by the bulk, so this observation could be considered as a confirmation of the existence of a B flux from the surface towards the bulk. Volatilization of B is unlikely as HfB2 and B are  Journal of the European Ceramic Society 38 (2018) 47-56  stable in low vacuum atmospheres (10−2 Torr). Fig. 5 shows an SEM micrograph (top left) of the microstructure of HS20 after deformation at 2000 °C. In addition, different composition maps are shown as well as a line scan analysis. Moreover, B and Hf are detected in some areas from the original SiC secondary phase (dark contrast) in addition to Si and C being present in the HfB2 matrix, which is also shown by the line analysis. A more in depth analysis by TEM reveals the formation of different phases as shown in Fig. 6. Different particles in the ternary Hf-Si-B system were detected by EDS after deformation at 2000 °C (Fig. 6a) and 1500 °C (Fig. 6b), suggesting bulk diffusion of B to produce a reaction forming new phases with different stoichiometries in the Hf-B-Si system. In addition, SiBn particles were detected at 2000 °C (Fig. 6a). SiBn particles with C-solid solubility and n as high as 32 have been observed previously [49]. Stoichiometry of phases determined by the semi-quantitative EDS used here is only approximate due to the difficulty of quantifying B by EDS. Some EDS spectra are shown in Fig. 7 as an example. After deformation at 1500 °C (Fig. 6b), 2 phases with brighter contrast in the Hf-Si-B-C and Hf-B-Si systems were detected by TEM. Dislocations are also present in the HfSi-B phase, in the magnified area indicated by an arrow, which is consistent with the observed stress exponent of 3 at 1500 °C. Cavitation damage, highlighted with circles in Fig. 6b, is also observed in samples deformed at 1500 °C, which could be responsible for the observed softening in the deformation curve. In addition, no signs of dislocations were observed at 1900 °C or 2000 °C consistent with either pure diffusional mechanisms or grain boundary sliding for micron-sized grains and measured stress exponents of 1 [36,37]. Talmy et al. [27] suggested grain boundary sliding as the deformation mechanism after studying the tensile side on 20 vol.%. SiC-reinforced ZrB2 after flexure creep testing and observing the crack propagation. In addition, Bird et al. determined the activation of grain boundary sliding above 1600 °C, whereas pure diffusional mechanisms were active below 1600 °C [29]. The latter conclusion was supported by the observation of cavities being responsible for accommodating deformation. However, none of the previous studies carried out any detailed microstructural characterization by TEM. For example, cavities are observed in this study at 1500 °C, in which dislocations are observed (Fig. 6b) and a stress exponent of 3 is measured that is in agreement with a power law creep mechanism instead of grain boundary sliding. Typically, in ceramics with a grain size between 1 and 5 μm, a stress exponent between 2-1 is related to grain boundary sliding under compression [36,37]. However, this is true for ceramics without any phase evolution during  Fig. 5. SEM micrograph of HS20 after deformation at 2000 °C (top left) with the corresponding elemental mapping of B, Hf, C and Si. Also, a line analysis through the arrow in the  micrograph is shown in the top right graph.  51  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  Fig. 6. a) BF TEM image of HS20 after deformation at 2000 °C. b) BF TEM image of HS20 after deformation at 1500 °C. Elements with concentrations below 5 at.% are considered as  impurities and written as subscripts in the chemical  formula.  Fig. 7. EDS spectra of; a) BHf,Si, b) (Si3B3O2)Hf,Mg and c) (Hf3Si2B)Mg.  deformation, such as YTZP, Al2O3 or SiC among others [37]. Therefore, the possible activation of pure diffusional mechanisms in the temperature range 1900 °C-2000 °C cannot be ruled out due to phase evolution or formation of new phases in Hf-B-Si/Hf-B-Si-C systems during deformation. Furthermore, even in the case that grain boundary sliding is the active deformation mechanism, it requires the parallel activation of pure diffusional phenomena to produce the observed phase evolution. In a recent study, Bird et al. confirmed the activation of grain boundary sliding during deformation of ZrB2-20 vol.% SiC (ZS20) as indentation misalignments were observed after deformation as a consequence of grain boundary sliding, i.e. grain rotation and translation [31]. However, the activation energies reported were 364 and 639 kJ/ mol for temperatures below 1600 °C or above 1600 °C [29], respectively, which are far higher than the ones reported by Talmy et al. [27], Gangirredy et al. [26] and this study, which are all in agreement. In fact, the creep rates reported by Bird et al. for ZS20 at 1800 °C are comparable to those reported for HS20 at 2000 °C in this study and similar to those reported for ZS30 at 1900 °C. This suggests that something is not being considered in Bird. et al. studies as their ZS20 should be more creep resistant than ZS30. It is noticeable in Fig. 15 of Bird et al. [29] that creep strain rate increases around 2 orders of magnitude from 1500 °C to 1550 °C under 100 MPa, while the difference in the creep strain rate at 100 MPa over all other temperatures shown is below an order of magnitude for a temperature difference of 100 °C. In addition, the experimental procedure shows that the powders were milled with WC-Co milling balls, which produced a WC-Co contamination of 2.3 wt.% [29]. Co melting point is 1500 °C highlighting how a relatively small amount of contamination can have a big impact on the high temperature creep response, which was not considered in the original study [29] or further microstructural characterization studies [31,50]. B diffusion through HfB2 and SiC produces phase instabilities during plastic deformation at high temperature, as observed in this work, which is responsible for the lower creep resistance compared to pure  (Si3B3O2)Hf,Mg  SiC in Fig. 2b. Fig. 7 shows different phases with different stoichiometries observed and characterized by EDS, such as BHf,Si (Fig. 7a), (Fig. 7b) and (Hf3Si2B)Mg (Fig. 7c), in which the subscript elements are considered as impurities present as their content is below 5 at.%. A TPS material must be structurally stable with no phase instabilities. As a result we need to look for alternative UHTC systems from HS20. In addition, the fact that HfB is stable and could be formed according to the Hf-B phase diagram might explain the observed instabilities as some B could be released from HfB2, diffusing to SiC particles [51] and leaving remnant HfB. ZrB2-SiC ceramics might be more stable at such high temperatures since the Zr-B phase diagram does not show the formation of ZrB [48] as discussed previously by Portnoi and Romashov who concluded that if it forms it must be stabilized by the presence of some impurities, such as O or N [52]. Note that the equilibrium conditions of phase diagrams at high temperature do not correspond with our experimental conditions in the presence of stresses and subsequent deformation at high temperature. This might lead, for example, to effects such as reported by Seifert and Aldinger [53] who found SiC was stable in the presence of B4C in the Si-B-C ternary system at 2500 K [48]. However, the presence of Hf is not addressed in their study [53] which could influence the stability of SiC in the Hf-Si-C-B quaternary system as Si diffuses towards HfB2 (Fig. 5). A comparison of HS20 microstructural evolution with ZrB2-SiC UHTCs is not possible due to the lack of microstructural characterization of studies published to date carried out by TEM. Also, HfB exists in the binary Hf-B diagram while ZrB does not exist in the binary Zr-B diagram [48,51,52], which could result in different phase stabilities at high temperatures in Zr-Si-C-B and Hf-Si-C-B quaternary phase diagrams. For example, Gangireddy et al. [26] suggested symmetric plastic behavior according to the TEM study of Jayaseelan et al. [30], in which they did not observe any amorphous phase in the ZrB2-SiC interphase. In our study, some borosilicate particles with minor Hf content were detected but only in samples which had been deformed at 2000 °C (Fig. 6), which suggests that the plastic response might not be symmetric, in agreement with Bird. et al. study [31]. Also, the plastic  52  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  response from a ZrB2-30 vol.% SiC is similar to the HS20 studied here as there is only a factor of 2-3 difference in the strain rate for the steady states of the plastic deformation at the same stress level and same temperature. In addition, the fact that the activation energies agree well with those in the ZrB2-SiC studied by Talmy et al. [27], suggests that similar mechanism to accommodate the deformation might be active (grain boundary diffusion). Neuman et al. [23] reported formation of liquid B-O-C-N inclusions in the only study including a TEM image after deformation from 1800 °C − 2300 °C, which are responsible for the strength degradation in this temperature range. However, microstructures after strength testing cannot be compared with those developed after compressive creep as the strength tests by Neuman et al. were measured under flexure and produced elastic failure, without there being time for diffusion to play a role in the microstructure development. Furthermore, an increase in the SiC cluster size was observed when testing at or above 1800 °C through post mortem SEM only [23]. Therefore, it is important that future studies give improved microstructural characterization to enable a better understanding of the phenomena involved in high temperature deformation. Melendez-Martinez et al. [44] showed; (i) ZrB2 at 1500 °C is 200 times more creep resistant than HS30 and (ii) the deformation mechanism is grain boundary sliding and that the deformation of individual ZrB2 grains could be neglected as well as any contribution to deformation from ZrB2-ZrB2 grain boundaries. Therefore, grain boundaries of secondary phases with ZrB2 and with the secondary phases are responsible for the macroscopic deformation and link the B diffusion to form new secondary phases in the Hf-B-Si and Hf-B-Si-C systems as the existing accommodation process of the deformation mechanism. However, it is difficult to assess how the formation of these secondary phases start and how each phase contributes to macroscopic deformation. Future studies at low deformations will try to answer how instabilities are triggered and how deformation is controlled in the initial stage of plastic deformation. In summary, our study suggests that HS20 behaves at high temperature as follows; (i) a power law creep with thermal activation process accommodated by B lattice diffusion (n=3 and Q=242 kJ/mol) in the temperature range 1500-1800 °C, consistent with the observation of dislocations by TEM, (ii) grain boundary sliding/pure diffusional mechanism accommodated by B grain boundary diffusion (n=1 and Q=334 kJ/mol), consistent with the absence of dislocations and the generation of new phases in the HfB-Si and Hf-B-Si-C systems.  3.3. HS20La high temperature mechanical properties  2 wt.% addition of La to HS20 (HS20La) helps reduce the internal stresses produced during cooling in the SPS without altering mechanical properties at RT, such as strength and toughness [19]. In addition, La addition stabilizes the thermal conductivity in the temperature range from 1500 °C to 1900 °C due to a compensation of CTE mismatch between SiC and HfB2 as La2O3 is located next to SiC particles [6], which could be an advantage for designers when calculating UHTCs performance [3]. Moreover, La2O3 addition stabilizes the long-term response during oxidation at 1500 and 1600 °C due to the formation of Hf-O-C particles which slow down growth of the oxide protective layer [11]. Therefore, unlike HS20 HS20La could be a good candidate for hypersonic applications. However, Fig. 8 reveals that the plastic response is far from being satisfactory for several reasons; (i) There is a sudden softening after the end of elastic deformation suggesting development of structural damage with the plastic deformation, (ii) the lack of a true deformation stationary as softening is observed under any of the conditions studied, and (iii) HS20La is even less creep resistant than HS20 (Fig. 8b). There was no attempt to calculate activation energies as no true stationary state was observed under any condition. In addition, the stress exponents given in Fig. 8 are only indicative to enable a quick comparison with HS20 but they were not calculated with the intention of suggesting a deformation mechanism.  La2O3 addition does not act as a high temperature structural reinforcement. In addition, powders were observed immediately around the sample after testing at 1800 °C, a smaller amount (< 10 mg) of powder was observed at 1900 °C but no trace of powder was observed at 2000 °C, which could be related to the development of either porosity next to the external surface or a chemical reaction during deformation. C content of the powders found at 1800 °C was 1.70 ± 0.02 wt.%. However, the amounts of powder collected at 1900 °C were not enough for C content analysis. Considering that C content for HS20La is 2.01 wt.%, the latter finding suggests that SiC might segregate in Si and free C at 1800 °C. Moreover, the higher the temperature, the lower the porosity detected at the external surface as diffusion phenomena are more active balancing the presence of any damaging mechanisms such as cavitation or reaction (Fig. 9). For example, Fig. 9a shows the cross section after compression testing at 1800 °C, in which the highest degree of porosity is observed and is reduced as temperature increases to 1900 °C (Fig. 9b) and 2000 °C (Fig. 9c). In a previous study, La2O3 particles were observed to be located next to SiC particles [19] which could have an influence on the SiC disociation as this phenomena was not observed in HS20. However, after deformation at 2000 °C there was no O in what were originally La2O3 particles as shown in Fig. 10. The remnant La2O3 particle contains La-Si-Hf-B after deformation, which suggests that O might react with C from SiC particles. Also, the stoichiometry of the SiBn particles is different in HS20 than HS20La: the maximum n in HS20 is 2 whereas it is 12 for HS20La. In addition, SiBn particles with n > 2 contain minor amounts of La, indicating that La plays a role in the formation of new SiBn phases during deformation, which are not formed otherwise. SiBn as a solid solution with n as high as 32 has been reported [49]. Also, some of the SiBn decomposes through a peritectic reaction involving liquid formation, e.g. SiB6 (observed in Fig. 10), which decomposes to a liquid plus SiBn with n > 6. The decomposition temperature of SiB6 is 1850 °C but how Hf and La could affect the stability of SiB6 is unknown. Our study indicates that if decomposition reaction occurs, it has to be above 1900 °C as a stress exponent of 3 could be a clear indication of dissolution-precipitation deformation mechanism in ceramics where a liquid is present at the grain boundaries. This is also indicated by the presence of rounded grains in Fig. 10 on SiBn particles with n > 6. Disociation/Decomposition of SiC or SiBn phase could be responsible for the mechanical resistance degradation observed in the stress-deformation curves (Fig. 8). However, to clarify which phase is responsible for the mechanical response degradation, the initial stage of deformation should be studied in detail, which is beyond the scope of this work and will be the subject of future studies. Nonetheless, the poor creep resistance shown by HS20La highlights an unsatisfactory behavior for hypersonic re-entry involving temperatures in excess of 1800 °C. Besides, the use of RE-oxides or any oxide is not recommended according to these results in spite of having the ability to either develop dense-oxidation-resistant RE-zirconates/hafnates layers or stabilize oxidation formation through the presence of Zr-O-C/Hf-O-C particles. Assuming the conditions of temperature and stresses described above, 1800 °C and 400 MPa, and the results from this study (Fig. 8), we suggest (i) HS20 is deforming at 3 × 10−4 s−1 and (ii) HS20La at 10−3 s−1. In addition, plastic deformation at 1800 °C starts at around 4% and 3% for HS20 and HS20La (Fig. 1 and 8), so that at stresses below 400 MPa there will be no dimension changes for 130 s and 30 s respectively. Moreover, 10% deformation will be reached in 330 s and 100 s respectively. In a sharp leading edge component the deformation should not go above 10% under any circumstances due to the negative impact on flight aerodynamics. Therefore, considering that an atmospheric re-entry lasts around 10 min being at least half of the time under the hardest conditions, i.e. 1800C and 400 MPa in this discussion, only HS20 would be able to withstand it with a deformation below 10% (9% approximately) while HS20La would be far above 20% deformation, mechanically failing before the end of the re-entry. Ideally, the material should not exceed elastic deformation to be reusable from  53  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  Fig. 8. Stress versus strain for HS20La at; a) 1800 °C, b) 1900 °C and c) 2000 °C. Strain rates for each deformation state are shown. d) Strain rate versus stress for the deformation  stationaries of HS20 and HS20La.  Fig. 9. Cross section image of HS20La after deformation at; a) 1800 °C, b) 1900 °C and c) 2000 °C.  Fig. 10. BF TEM picture of HS20La after deformation at 2000 °C with an EDS spectra for the phase with a *.  54  \\x0c', 'E. Zapata-Solvas et al.  Journal of the European Ceramic Society 38 (2018) 47-56  a structural point of view although those materials exceeding the elastic deformation developing a certain shape change such as HS20 might be suitable for single use. Considering other UHTCs studied, ZrB2-30 vol.% SiC deforms at 5 × 10−4 s−1, which means that it is slightly out of the secure range defined in this study as a 10% deformation and the elastic threshold (4%) will be reached in 200 s and 80s, respectively. ZrB2-20 vol.% SiC samples tested under plasma wind tunneling testing conditions showed the tip of a sharp leading edge was slightly curved and deformed afterwards even if testing times were no longer than 50 s [54]. Note that 1800 °C and 400 MPa are mild re-entry conditions and ideally materials capable of withstanding re-entries with temperatures in excess of 2000 °C are needed in the long term. For example, none of the MeB2-based UHTCs analyzed in either this study or previous studies could survive without deforming less than 10% at 2000 °C under 200 MPa, which highlights a need for new UHTCs and approaches. In addition, there is a lack of studies about deformation mechanisms in the UHTCs community in the temperature range 1800-2000 °C and in the strain range rate of interest (10−5-10−4 s−1). Therefore, it is difficult to make an assessment about structural performance of UHTCs under realistic mechanical stresses at temperatures of interest. Nonetheless, the data reported till now indicate that alternatives should be explored. Furthermore, the creep resistance of MeB2 UHTCs is completely unknown and it should be studied before the UHTCs community can understand how these materials might be structurally reinforced.  4. Conclusions  High temperature compressive creep of HS20 and HS20La UHTCs up to 2000 °C reveals that HS20 might withstand the proposed conditions for a single use in hypersonic reentries but materials with higher creep resistance should be developed. RE oxide additions to HS20 degraded the creep resistance compared to HS20. Therefore, it is necessary to explore other systems as well as understanding the creep response of monolithic materials to indicate how they could be structurally reinforced.  Acknowledgments  The Authors’ acknowledge Prof. Mike Reece, Nanoforce Technology Ltd., Queen Mary, University of London, UK for providing access to the Spark Plasma Sintering facility. EZS also acknowledges support through a contract from the JAE-DOC program of CSIC, Spain, co-funded by the FSE to carry out this research project (2012-2015). DGG, ADR and EZS acknowledge the project MAT2015-71411-R from MINECO (Spain). 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},{
  "_id": 82,
  "PDF": "High temperature oxidation of two- and three-dimensional hafnium carbide and silicon carbide coatings.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 34 (2014) 879-887  Original Article  High temperature oxidation of twoand three-dimensional hafnium carbide and silicon carbide coatings  C. Verdon a,∗  , O. Szwedek a , A. Allemand a,b , S. Jacques a , Y. Le Petitcorps a , P. David b  a Université Bordeaux, LCTS, UMR5801, F-33600 Pessac, France b CEA, DAM, Le Ripault, F-37260 Monts, France  Received 26 April 2013; received in revised form 9 October 2013; accepted 15 October 2013  Available online 20 November 2013  Abstract  The difficulty  in using C/C composites as structural components above 2000 C  in an oxidizing atmosphere  is  their poor  lifetime. The solution proposed here consisted in combining two refractory carbides, hafnium and silicon carbides, in coating with a complex architecture, named a three dimensional coating, over a C/C substrate. Such a coating protects the C/C composite at 2000 C under air. The oxidation of the coating leads to the formation of a SixOyHfz hafnium-containing silicate  liquid, combined with HfO2(s) . This  liquid  limits oxygen diffusion more  than pure SiO2 does, so it is a better protection against oxidation. Furthermore, HfO2(s) acts as a frame holding SixOyHfz in place. From these results, an oxidation mechanism is proposed and discussed. © 2013 Elsevier Ltd. All rights reserved.        Keywords: HTC; Oxidation; Carbides; Refractories; Hafnium  1.   Introduction     C/C composites  (carbon fibers embedded  in carbon matrices) exhibit a poor lifetime under an oxidizing atmosphere; the carbon fiber starts  to be oxidized under air at a  temperature of 400 C.1 In order  to protect composites from oxidation, many coatings made of Ultra High Temperature Ceramics (UHTCs) have been tested. It is known that they can resist to extreme heat flux and high mechanical stresses.2 For example, Clougherty et al. have studied oxidation of diborides HfB2 , ZrB2 TiB2 and a mixture of HfB2-SiC.3 The oxidation mechanism model of these coatings has been studied by Parthasarathy et al.4 who proposed a mechanistic model that simulates the oxidation behavior of  the diborides  in  the  temperature  range of 1000-1800 C. Among  the  studied carbide-based coatings, an HfC/SiC dual layer developed by Wunder et al. has shown promising oxidation protection for PyC-coated C/C composites at temperatures up  to 1450 C.5 Besides, Baklanova et al. highlighted  the fact that HfC/SiC-coated carbon fibers exhibit a higher oxidation C than the initial HfC-coated carbon fibers.6 resistance at 2000           ∗  Corresponding author. Tel.: +33 556 84 47 30; fax: +33 556 84 12 25.  E-mail address: verdon@lcts.u-bordeaux1.fr (C. Verdon).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.10.019                 Wang et al. have determined that SiC/HfC/SiC is an interesting coating  to protect carbon/carbon composites with a mass  loss of only 2.3% after oxidation at 1500 C. They have also characterized a stable glassy phase composed of HfSiO4 which could protect C/C composites.7 In the present work, two carbides were selected. SiC has been used in many works; it allows protection against oxidation until C by forming a SiO2 layer.8 It is the best diffusion barrier 1650 C.1 But, above 1650 against oxidation below 1400 C under air  the oxidation becomes active and all  the protection  is  lost. In spite of a high melting point (3890 C), HfC is actively oxidized at low temperature, from 400 C to 500 C, forming porous HfO2 . Its oxidation behavior has been studied by Shimada et al.9 Although not protective,  this oxide has  the advantage of being refractory with a melting point of 2810 C and has a lower vapor pressure than SiO2 .4 The aim of the present study was to combine these two carbides to find a synergetic effect: a refractory coating that allows a good oxidation resistance. Three kinds of samples, (i) a monolithic HfC/SiC samples type consisting of coated and sintered powder, and two C/C composites with (HfC/SiC)n multilayer coating  types,  (ii) 2D and  (iii) 3D, were prepared and patented.10,11 In the 2D (HfC/SiC)n multilayer coatings delamination can occurs and layers could slide over each other during oxidation at high  temperature. The 3D (HfC/SiC)n multilayers                 \\x0c', '880   Table 1  Sample characteristics overview.  Sample kind   Sample composition  Way of synthesis  C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887  Sample A  Monolithic   HfC/SiC   FBCVD + SPS  Sample B   2D (10 layers)   C/C substrate + (HfC/SiC)5 CVD   Sample C  3D + 2D (10 layers)  C/C substrate + (HfC/SiC)5 CVD  )  C  °  (  e  r  u  t  a  r  e  p  m  e  T  2500  2000  1500  1000  500  0  0  re (°C)  temperatu  pressure  (M   Pa)  5  10  15  20  25  30  Time (min  )  p  r  e  s s  u  r  e  (  M  P  a  )  80  70  60  50  40  30  20  10  0  35  1st layer SiC  2nd layer HfC 3rd layer SiC  4th layer HfC  5th layer SiC  Fig. 2. Structure scheme of the 5 first layers of sample B corresponding to 2D  structure coating.  Fig. 1. Pressure and temperature cycles of spark plasma sintering.  1st  layer SiC whiskers  2nd layer HfC  3rd layer SiC  comprise SiC whiskers deposited as a first  layer. As explained by Chu et al., SiC nanowires (or whiskers in our case) increase the mechanical properties of a coating such as hardness, elastic module and  fracture  toughness and  improve  the oxidation protection of the sample.12,13 These improvements are due to a mechanism  including whiskers pullout, micro-crack deflection and better interface interaction, whiskers acting as attachments of sub-layers. Then samples were characterized after high temperature oxidation. This characterization of the oxidized samples allowed us to propose an oxidation mechanism.  2. Experimental procedure  2.1. Monolithic sample  sample A,  The monolithic  freestanding  sample  (named  Table 1) was prepared by following two steps. HfC powder with a d50 = 35  \\u242em, was first coated with 1  \\u242em of SiC by fluidized bed chemical vapor deposition  (FBCVD) by Lifco  Industrie  (France). Then  this core-shell powder was sintered by spark plasma sintering (SPS). Sintering parameters were:  temperature of 1950 C, pressure of 75 MPa and dwell time of 5 min, following cycles presented in Fig. 1. The machine was a Dr Sinter 2080 from Syntex (Japan). The final sample was a cylinder with a diameter of 15 mm and a thickness of 6 mm.     2.2. Multilayer coating made by CVD  The other kinds of samples (sample B and sample C, Table 1) consisted in (SiC/HfC)n multilayer coatings over C/C substrates prepared by  low pressure chemical vapor deposition (CVD) at 1000 C and 5 kPa.11 Samples B and C have  the same size as     HfC 4th layer  SiC 5th layer  Fig. 3. Structure scheme of the 5 first layers of sample C corresponding to 3D  structure coating.     sample A. The device used was composed of a hot-wall CVD reactor and a chlorinating device. HCl(g) reacted at 700 C with metallic Hf(s) to form HfCl4(g) in  the chlorinating device. The hafnium metal was an electrolytic grade Hf(s) supplied by Areva. HfCl4(g) precursor was simultaneously injected with argon as a carrier gas and reacts with CH4(g) and H2(g) in the CVD reactor to give HfC coating. The SiC layers were classically made from methyltrichlorosilane and hydrogen. The MTS was supplied by Sigma-Aldrich and its purity was superior to 97%. Sample B consisted of a classical bi-dimensional (2D) coating made of  ten alternated  layers of SiC and HfC, according to Kaplan’s patent.14 A structure scheme, corresponding to the five first  layers,  is presented  in Fig. 2. The five  last  layers are identical to the five first ones. Sample C consisted of five alternated  layers of HfC and SiC with a  three-dimensional  (3D) arrangement and another five alternated  layers of HfC and SiC with a  two-dimensional arrangement, according to our patent.11 In this 3D structure, the          \\x0c', 'C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887   881  HfC SiC Porosities  Fig. 4. Temperature profile applied for the oxidation test.  Fig. 6. Polished section of monolithic sample observed by SEM.  20 µm 20 µm  first  layer  is made of SiC whiskers  instead of a classical dense continuous SiC layer. A structure scheme corresponding to the five first  layers  formation of  the 3D structure  is presented  in Fig. 3. The last five layers with a 2D arrangement are identical to the five first ones of sample B.  2.3. Oxidation tests  The samples were oxidized under air  in an arc  image  furnace at the CEA/Cesta center (Bordeaux, France). The furnace is composed of six Xenon arc  lamps. These  lamps are placed in  the first focus of elliptical reflectors  in order  to concentrate emitted radiation at the second focus. The second focus of each reflector coincides in one place, corresponding to sample position. It allows adding six  lamps energy on  the surface sample. The temperature of the center of the samples measured by a thermocouple reached 1600 C in 50 s and finally 2000 C in 200 s (Fig. 4). Then samples are kept 5 s at 2000 C. Seven samples were oxidized. This duration and  temperature of oxidation  test correspond  to sample B coating destruction. The center of  the surface sample was the most affected area due to a temperature gradient between the center and the side of the surface sample.           2.4. Characterization  The crystal  structure of  the  sample  surface was  identified by X-ray diffraction (XRD) with a Diffractometer D8 Advance  Bruker  (CuK␣). Before  and  after oxidation  test,  specimens were mounted  in  resin, cut and polished  in order  to carry out morphological observation by  scanning  electron microscopy (SEM) with a SEM FEI Quanta 400 FEG using backscattered electron  (BSE) and  secondary electron  (SE) detectors under an accelerating voltage  less  than 20 kV. Qualitative chemical analyses were also performed with  the microscope by energy dispersive X-ray spectroscopy (EDS). The composition of each layer after oxidation was determined more precisely by electron probe microanalysis (EPMA) with an EPMA CAMECA SX100. Raman spectroscopy analyses were realized with a Horiba JobinYvon, Labram HR spectrometer (λHe-Ne = 632.8 nm).  3. Results  3.1. Monolithic sample  Fig. 5(a) presents an HfC particle before FBCVD. Fig. 5(b) corresponds  to a back scattered electron picture of HfC coated powder,  coating  is uniform  almost  every particle  is  coated. Coated particles are gray and uncoated particles are white. During spark plasma sintering bonding between particles of core-shell powder is created. Particles are sintered together into a monolithic sample.15 Fig. 6 presents a polish cross section of a monolithic sample before oxidation (sample A). A homogeneous SiC layer is visible around particles.  (a)  (b)  Uncoated particle  HfC  HfC particle  10 µm 10 µm  Coated particle  HfC  300 µm 300 µm  Fig. 5.   (a) Morphology of the HfC particle before FBCVD. (b) Back scattered electron of HfC powder coated by SiC, bright particles are uncoated.      \\x0c', '882   C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887  )  .  u  .  a  (  HfO2 (JCPDS 04-005-4474)  y  t  i  s  n  e  t  n  I  13  20  30  40  50  60  70  80  90  22θθ  Fig. 7. XRD pattern of sample A surface oxidized.  Sample A is the most convenient to investigate the oxidation behavior because its important thickness allows all the oxidation steps  to be separated. The sample has not been affected  in  the bulk by oxidation. After oxidation, following the path exhibited in Section 2.3, the sample is studied by XRD. The analysis of  the sample surface gives a pattern  typical of crystallized oxide HfO2(s) corresponding to the JCPDS card N 04-005-4474 (Fig. 7). HfO2 is  the only detected phase. The pattern of the untouched material is not visible under these thick oxide seals.     Table 2  Average composition of sample areas.  Average composition (%at)  Element   Area 3   Area 2   Area 1  Hf   Si   O   C   21   15   57   7   28   0.5  41   30.5   33  7  2  58  \\u242em   The sample is composed of 3 areas observed by scanning electron microscopy in Fig. 8(a) Area 1 is undamaged and exhibits a thickness of 59 mm (Fig. 8b). The initial thickness is 60 mm. It  is composed of HfC grains with SiC at HfC grains boundary, no change of grains sizes can be highlighted compared  to initial particles. A  few porosities are visible. Area 2  is highly porous and  is around 100  thick (Fig. 8). Silicon could not be found  in Area 2. Area 3, at  the outer surface,  is composed of an oxide skeleton in a glassy continuous phase and is around 250  \\u242em thick (Fig. 8d). EDS analyses  show  that  the glassy phase consists mainly of silicon (higher  than 30% atomic) with oxygen (higher  than 60% atomic) and some hafnium (less than 5% atomic) according  to  the formula SixOyHfz . The oxide skeleton  is composed of hafnium and oxygen  in proportion of HfO2 , confirmed by EPMA. The oxidation area total thickness is around 350  \\u242em. Fig. 9 corresponds  to component quantities observed during EPMA analysis. Table 2 lists average quantities of elements detected  in each area. The hafnium atomic composition determined  by  EPMA  is  quasi-constant  on  the  entire  sample  (a)  HfC(s)  (b)  SiC(s)  Area 3  Area 2  Area 1  HfO2(s)  (c)  EPMA analysis  400µm 400µm  Porosities filled by resin  50µm 50µm  HfO2(s)  (d)  Si  xOyHfz  50µm 50µm  20µm 20µm  Fig. 8. SE SEM observation of sample A after oxidation at 2000     C and quenched under air: overview image with the 3 areas that compose this sample (a), example  of enlarged view of area 1 composed of 2 carbides (b), example of enlarged view of area 2 composed of HfO2(s) (c), example of enlarged view of area 3 composed of two oxides (HfO2 and SixOy Hfz ) (d).                      \\x0c', 'C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887   883  Area  3  Area  2  Area  1  (a)  HfO2(s)  recrystallized  HfO2(s)  (b)  HfO2(s)  recrystallized  Porosities filled by resin  Particle boundary  Hf  Si  O  C  50 µm 50 µm  5 5 µm µm  Fig. 11. View of area 3 containing dendrites made by   recrystallization of a  Hf-containing compound.   (a) Corresponds   to an axial view of dendrites.   (b)  Corresponds to a perpendicular view of dendrites.  50  0 100  50  0 100  %At  50  0 100  50  0  0            50          100          150          200          250  µm  Fig. 9. EPMA analyses of sample A after oxidation.  Area  3  Area  2  100 µm 100 µm  Fig. 10. SE SEM observation of porosities at the surface of the oxidized sample  A.  (Figs. 8 and 9 and Table 2) aside from holes and particle boundary. Silicon is present in areas 1 and 3; area 2 is silicon-free. The atomic percent of carbon decreases from area 1 to area 3. In area 2, carbon peaks correspond to porosity filled with carbon resin. In area 3, carbon  is detected no more.  In  reverse,  the oxygen atomic content  increases  from  the  inside  to  the outside of  the sample (Figs. 8 and 9 and Table 2). The compositions of the three areas can be deduced from this analysis. Area 1  is a non affected part composed of HfC/SiC grains. Area 2 is made of porous HfO2(s) . Area 3 is composed of HfO2(s) skeleton combined with amorphous SixOyHfz (Fig. 8 d). The SixOyHfz glass is, of course, not detected by XRD because of its amorphous structure (Fig. 7). The microstructure of  the oxidized  sample A exhibits  the presence of holes in area 3 which results from gas formation in large quantities (Fig. 10). A particularity of Area 3 is Hf-containing recrystallized dendritic zones, more visible  in  important glassy phase volume (Fig. 11). Raman spectroscopy analyses results are presented in Fig. 12. Spectra of  the HfO2(s) skeleton were first acquired  (Fig. 12a) the spectrum exposed by Tkachev et al.16 As and compared  to  expected, these spectra are typical of HfO2 crystals. Then, spectra of the parts composed of the glassy phase and dendrites were acquired (Fig. 12b). The Raman peaks are identical in both cases.  Fig. 12. Raman spectrum. (a) Curve corresponding   to glassy phase with den drites analyses.   (b) Curve corresponding   to HfO2(s)  in area 3   formed during  oxidation tests.  HfC  SiC  Substrat  e  resin  10 µm 10 µm  Fig. 13. SE SEM image of a sample B-type coating, made by CVD. This sample  consists of a   two-dimensional multilayer coating on C/C substrate   (only five  layers are shown on this picture).  Consequently, the dendrites can be identified as HfO2(s) within a glassy phase.  3.2. CVD coatings on C/C composites  3.2.1. Sample B  Fig. 13 presents a sample B-type coating over a C/C substrate with  three  layers of SiC  in dark gray and  two  layers of HfC  in light gray. Unlike Fig. 13, sample B has actually a  total of  ten alternated layers of HfC and SiC, the coating is 20  \\u242em thick (not shown here). Fig. 14 presents the microstructure after oxidation at 2000 C under air after 200 s of sample B (same oxidation step as sample A and C). Neither residual carbide multilayer coating nor oxide                 \\x0c', '884   C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887  resin  Cracks = Oxidized  area Coating is no  more visible  Damaged substrate  C/C  Undamaged C/C substrate  50 µm  Fig. 14. SE SEM picture of a sample   initially coated with a   two-dimensional     C. No coating   is visible  coating,   like sample B, after oxidation   in air at 2000  anymore; it delaminated.  3D  structure  2D  structure  Outer  surface  30 µm  30 µm  HfC  SiC  SiC  Fig. 15. SE SEM images of sample C before oxidation.  layers are visible. The C/C substrate is naked and only slightly damaged at the surface. It is likely that an oxide scale has protected this sample. But a delamination of the superficial oxidized layer occurred.  3.2.2. Sample C  \\u242em   Fig. 15 presents  an  example of  sample C before oxidation with  its 10 deposited  layers. The coating  is 40  thick. Two parts are visible,  the five first  layers  forming  the  threedimensional structure where  the sub-layers are concentrically arranged around SiC whiskers. This structure was detailed  in Section 2.2. Over  this structure a  two-dimensional part made of five  layers  is visible.11 The first  layer of sample C  is composed of SiC whiskers. Consequently when  the next  layers are deposited  it  is possible  that  some pores are  formed early by closing open space between close whiskers without filling completely  the spaces between  the whiskers. The effect could be negative but  it  is difficult  to assess because pores also appear during oxidation of sample as a result of gas release. After oxidation (following  the path exhibited  in Section 2.3 oxidation  tests) of sample C, an oxidized coating  layer  is still observed on the C/C (Fig. 16). Fig. 16(a) presents the center of the sample surface, where  the coating has been more severely oxidized. Wide cracks have been filled with mounting resin. At the outer surface, a SixOyHfz glass  is evidenced by chemical analysis, as in sample A. Holes created by gas released are also observed. The  initial  two-dimensional  (SiC/HfC)n multilayer  coatings have been replaced by some HfO2(s) structure broken into disorganized pieces. SiC seems  to have disappeared; only some residual parts are visible in Fig. 16. The three-dimensional and undamaged structure of the coating is still visible close to the substrate. In place of the initial three-dimensional coating close to the substrate, the oxide mixture of HfO2(s) and SixOyHfz is also found but with a more refined microstructure. The carbon/carbon substrate has begun to be locally damaged. Fig. 16(b) presents the edge of the sample, where the coating is less damaged. The different behavior between the side and the center of  the sample surface can be explained by  the  temperature gradient. Under  the outer SixOyHfz glass  layer, evidenced by chemical analysis, HfO2(s) sub-layers are found separated by wide cracks filled with mounting resin. This feature evidences that SiC layers have undergone an active oxidation. The free created spaces become delamination areas between the sub-layers. Although cracked  in  some places,  the oxide  layers are more continuous  than  in Fig. 16(a) and  retain  the original shape of the  two-dimensional HfC  sub-layers  that  they  replace. Some residual SiC sub-layers are still present. Close to the substrate, the  initial  three-dimensional multilayer carbide coating  is still observed with  its undamaged structure. The  inner part of  the coating is thus non-oxidized, SiC and HfC layers being intact. From the studies of these samples, an oxidation mechanism can be proposed. The species found  in sample C are  the same as in sample A after oxidation. The hypothesis is that the mechanism of  the monolithic sample oxidation and  the one of  the three-dimensional coating oxidation are very similar and involve the same chemical reactions.  4. Discussion  4.1. Mechanism based on the monolithic sample behavior                    Pure HfC begins  to be oxidized above 500 C.9 But  in  this study, HfC is combined with SiC. From 500 C to 1500 C, SiC is oxidized  to SiO2 which  is supposed  to significantly protect HfC from oxidation because SiO2(s) is  the most effective oxyC.1 From room temperature to gen barrier coating below 1500 1600 C the oxidation temperature rises in less than 100 s which can be considered as very fast. Consequently, it is admitted that the SiO2 first layers produced by SiC oxidation still protect HfC from O2 which did not have the time to diffuse through the SiO2 layer. Above 1700 C,  in agreement with  the HfO2-SiO2 binary diagram (Fig. 17),17,18 the suggested mechanism of sample A oxidation, presented in Fig. 18, consists of 3 steps which, once initiated, can take place simultaneously. First, HfC and SiC are oxidized according to reactions R1/R2 and R4/R5 near the outer surface depending on the oxygen partial pressure. These reactions release gases (CO(g) , CO2(g) and SiO(g) ). At the surface, SiO(g) released from the inner part of the sample is oxidized according to R3 to form a condensed and liquid layer of SiO2(l) because of higher oxygen partial pressure.19 This liquid tends to seal the pores and cracks of the outer scale. Moreover CO(g) and CO2(g) create bubbles  in  the outer  liquid oxide  layer as shown  in Figs. 10 and 16(a)  (where holes are                \\x0c', 'C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887   885  a)  resin SixOyHfz Holes  HfO2(cr)  SiC Cracks filled with  resin Intact HfC  and SiC Substrate C/C  100µm  b)  40µm 40µm  Fig. 16. BSE SEM images of the center (a) and the side (far from the spot impact) (b) of sample C exposed to oxidation in air at 2000     C.  Fig. 18. Schematic representation of the oxidation mechanism of the monolithic  sample.  (Fig. 17)  to a composition of at  least 95 Mol% of SiO2 ,  that is completely  liquid at 2000 C. Thus,  the cooling was not fast enough  to completely prevent some HfO2 recrystallization  in dendritic shape. The presence of  this secondary HfO2 crystals further evidences that the oxide liquid formed at high temperature indeed contained hafnium and can be referred as SixOyHfz(l) . Hence, hafnium is always found by EDS analyses. But, contrary to Wang et al.7 the HfSiO4 glassy phase has not been observed here. This glassy phase has been formed during a longer oxidation time at lower temperature (1500 C and 66 h).  3/2O2 →   CO2(g) +   SiO(g)  R1 : SiC   +        R2 : SiC   +   1/2O2 →   CO(g) +   SiO(g)  R3 : SiO(g) +   1/2O2 →   SiO2(l)  R4 : HfC   +   2O2 →   HfO2(cr) +   CO2(g)  R5 : HfC   +   3/2O2 →   HfO2(cr) +   CO(g)  R6 :xSiO2(l) +  zHfO2(cr) →  liquid(SixOyHf z(l) )  From this, it is proposed that the effective protection against oxidation  is based on  the formation of  the SixOyHfz hafniumcontaining silicate liquid combined with the HfO2(cr) skeleton.  Fig. 17. HfO2 -SiO2 binary diagram.17     observed) and cracks in the HfO2(s) structure in the deeper and outer parts as shown in Fig. 16. The phenomena involved in the deeper part correspond to the formation of area 2 in the schematic representation (Fig. 18). Finally, SiO2(l) dissolves part of HfO2(s) to form a hafniumcontaining silicate  liquid, SixOyHfz(l) , according  to R6 and  the binary diagram of HfO2 -SiO2 that highlights  the coexistence at 2000 C of HfO2(s) with a silicon-rich liquid (SiO2 90%mol. ; HfO2 10%mol. ) (Fig. 17).17,18 The  liquid results  in a SixOyHfz hafnium-containing silicate glass after quenching leading to the formation of area 3 with the remaining HfO2(s) that was not dissolved by SiO2(l) and forms a skeleton. The presence of HfO2(s) dendrites has been evidenced in the glassy phase parts. The overall concentration of Hf  inside  this biphasic area  is  less  than 5% atomic.  It corresponds  in  the HfO2-SiO2 binary diagram      \\x0c', '886   C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887  η  As explained by Vogel, when a group 4 element, like hafnium, is added to SiO2(l) the tendency to phase separation and immiscibility rises and so  the viscosity  increases.20 Opeka et al.19 made the link between Vogel and the Stokes-Einstein relationship which shows that the oxygen diffusion coefficient through the liquid oxide is inversely proportional to the viscosity. Consequently, by adding Hafnium, the oxygen diffusion coefficient decreases and the next sub-layers are better protected from oxidation  than with pure SiO2(l) . These authors were  the first  to develop the concept of “phase separation as a controlling factor in the oxidation protection”. Additionally, the HfO2(s) skeleton acts as a frame holding the liquid SixOyHfz in place during oxidation and so helps to protect the next layers from further oxidation.  4.2. Extrapolation of the mechanism for CVD samples     The behavior under oxidation of  the monolithic sample A is used  as  a basis  that  can be  extrapolated  to  the oxidation mechanism of CVD coated samples B and C. In both samples B and C, the five exposed outer layers are twodimensional. First, the far outer layer, made of SiC, is oxidized below 1500 C  to form SiO2(l) according  to R1,  the oxidation of SiC being passive.8 When the temperature increases, the SiC oxidation becomes active and the SiO2(l) layer is ruptured by the produced gases. In  the same  time,  the oxygen diffuses  through the ruptured SiO2(l) scale and reaches  the deeper carbide sublayers. The oxidation mechanism should be then similar to the monolithic sample oxidation mechanism explained in Section 4.1. But as HfC and SiC have different coefficients of  thermal expansion, the temperature increase induces delamination and cracks between sub-layers that are facilitated by the rather flat geometry. Thus, contrary to sample A, neither HfO2(cr) skeleton within SiO2(l) nor SixOyHfz liquid phase can be  formed where  initial HfC and SiC are no more  in contact, each carbide phase being successively oxidized. The oxidation protection  is  thus less effective. Hence, in the case of sample B, as the whole coating has initially a two-dimensional structure, the C/C substrate is badly protected. In the case of sample C, the initial deeper three-dimensional structure  part  is  composed  of  concentric  and  interlocked carbide  sub-layers  that  give  a mechanical  grip  preventing delamination.20 Silicon and hafnium carbide phases being kept in contact with each other; the oxidation mechanism can occur as described above for the monolithic sample A and so protects more effectively the C/C substrate.  5. Conclusion  A new protection against oxidation at very high  temperature composed of hafnium and  silicon carbides was  realized by  two synthesis ways. Its behavior under oxidation was studied with an arc  image  furnace. After oxidation under air at 2000 C  the microstructure and composition of samples were characterized and a mechanism was proposed. Finally, the influence of  the coating  structure was  studied and  the advantage     of  the  three-dimensional  structure over  the  two-dimensional multilayer  structure was highlighted. The oxidation mechanism  took  into  account  six  reactions,  two  corresponding  to SiC  oxidation,  two  to HfC  oxidation,  one  to  an  oxidation of SiO(g) when  the oxygen partial pressure  increases  at  the sample  surface,  and one between HfO2(s) and SiO2(l) forming  a  hafnium-containing  silicate  glass  in  addition  to  the unreacted HfO2(s) ,  in  agreement with  the HfO2-SiO2 phase diagram. The  resistance against oxidation of samples mainly comes from  the SixOyHfz liquid phase which  limits  the oxygen diffusion better than pure SiO2 . In addition, the three-dimensional structure has three advantages. It allows the formation of HfO2(s) skeleton  that acts as a  frame holding  the SixOyHfz liquid  in place (i); consequently it cannot be easily removed by gas flux. The three-dimensional structure with initial interlocked carbide sub-layers  improves mechanical properties of  the coating12,13 (ii). Moreover,  it  increases  the exchange surface between  the chemical compounds (iii).  Conflict of interest statement  There is no conflict of interest.  Acknowledgments  This work was supported by  the “Alternative Energies and Atomic Energy Commission”  (CEA). We  thank M. Lahaye of  “the Center  for  characterization  of  advanced materials” (CeCAMA)  for  the characterization of  the  layers and Prof. F. Rebillat to give us tracks to follow.  References  1. McKee DW. Oxidation protection of carbon materials.   In: Thrower PA,  editor. Chemistry and Physics of Carbon, vol. 23. New York: Marcel Dekker;  1991. p. 173-232.  2. Paul A,   Jayaseelan DD, Venugopal   S,   Zapata-Solvas   E,   Binner   J,  Vaidhyanathan   ites   for hypersonic   B, Heaton A,  Brown  P,  Lee WE. UHTC  composapplications. Am Ceram Soc Bull 2012;91(1):22-8  http://hdl.handle.net/2134/9676  3. Clougherty EV, Pober RL, Kaufman L. Synthesis of oxidation resistant metal diboride composites. Trans Metall Soc AIME 1968;242:1077-82.  4. Parthasarathy TA, Rapp RA, Opeka M, Kerans RJ. A model  for  the 2007;55:5999-6010,  and TiB2 . Acta Mater   oxidation   of ZrB2 , HfB2 http://dx.doi.org/10.1016/j.actamat.2007.07.027.  5. Wunder V, Popovska N, Wegner A, Emig G, Arnold W. Multilayer coatings  on CFC composites for high-temperature applications. Surf Coat Technol 1998;100-101:329-32, http://dx.doi.org/10.1016/S0257-8972(97)00643-9.  6. Baklanova   NI,   Zima   TM,   Boronin   AI,   Kosheev   SV,   Titov   AT,  Isaeva NV, Grashenkov DV,   Solntesev   SS.   layer   coatings   for   carbon fibers.   Surf Coat Technol   Protective  ceramic multi2006;201:2313-9,  http://dx.doi.org/10.1016/j.surfcoat.2006.03.046.  7. Wang Y, Li H, Fu Q, Wu H, Yao D, Li H. SiC/HfC/SiC ablation resistant coating for carbon/carbon composites. Surf Coat Technol 2012;206:3883-7,  http://dx.doi.org/10.1016/j.surfcoat.2012.03.039.  8. Strife JR, Sbeehan JE. Ceramic coating for carbon-carbon composites. Am Ceram Soc Bull 1988;67(2):369-74.  9. Shimada   S,   Inagaki   M,   Matsui   K.   Oxidation   kinet ics   of   hafnium   carbide   in   the   temperature   range   of    \\x0c', 'C. Verdon et al. / Journal of the European Ceramic Society 34 (2014) 879-887   887  480     C   to   600     C.   J   Am   Ceram   Soc   1992;75(10):2671-8,  http://dx.doi.org/10.1111/j.1151-2916.1992.tb05487.x.  10. A. Allemand, L. Bianchi, Y. Le Petitcorps, O. Szwedek, Process for coat ing a part with an oxidation-protective coating, World patent application  N     WO2012055865.  11. A. Allemand, O. Szwedek, J.F. Epherre, Y. Le Petitcorps, Method for coat ing, using a chemical vapor deposition technique, a part with a coating for  protecting against oxidation, and coating and part, World patent application  N     WO2013076274.  12. Chu Y,   Fu Q,   Li H,   Shi X,   li K, Wen X,   Shang G.   Effect   of  SiC   nanowires   on   the mechanical  and  oxidation  protective  ability  of J Am Ceram Soc 2012;95(2):739-45,  SiC   coating   for C/C   composites.   http://dx.doi.org/10.1111/j.1551-2916.2011.04979.x.  13. Chu Y,   Li H,   Fu Q,   Shi X, Qi   L, Wei   B. Oxidation-protective  and mechanical properties of SiC nanowire-toughened Si-Mo-Cr  com2012;58:315-20,  posite   coating   for   C/C   composites.   Corros   Sci   http://dx.doi.org/10.1016/j.corsci.2012.01.003.  14. Kaplan R.B.   et   al. High   temperature   resistant   structure, US   Pat. N     5283109,1994.  15. Coble RL, Burke JE. Sintering ceramics. Prog Ceram Sci 1963:197-251.  16. Tkachev SN, Manghnani MH, Niilisk A, Aarik   J, Mändar H. Raman  and Brillouin   scattering   spectroscopy   studies of   atomic  layer-deposited 2005;61(10):2434-8,  ZrO2 and HfO2 http://dx.doi.org/10.1016/j.saa.2005.02.025.  thin   film.   Spectrochim Acta A   17. Levin EM, Robbins CR, McMurdie HF. Phase diagrams for ceramists. Am  Ceram Soc 1964.  18. Shin   D,   Arroyave   R,   Liu   Z.   Thermodynamic  model2006;30(4):375-86,  ing   of   the   Hf-Si-O   system.   Calphad   http://dx.doi.org/10.1016/j.calphad.2006.08.006.  19. Opeka MM,   Talmy   IG,   Zaykoski   JA.   Oxidation-based   materials  selection   for   2000     C + hypersonic   aerosurfaces:   theoretical  consid2004;39:5887-904,  eration   and   historical   experience.   J Mater   Sci   http://dx.doi.org/10.1023/B:JMSC.0000041686.21788.77.  20. Vogel W. 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},{
  "_id": 83,
  "PDF": "High temperature oxidation of Zr- and Hf-carbides Influence of matrix and sintering additive.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 33 (2013) 2867-2878  High temperature oxidation of Zrand Hf-carbides: Influence of matrix and sintering additive  Ludovic Charpentier a,∗  , Marianne Balat-Pichelin a , Diletta Sciti b , Laura Silvestroni b  a PROMES-CNRS, Laboratoire Procédés, Matériaux et Energie Solaire, 66120 Font-Romeu Odeillo, France b CNR-ISTEC, Institute of Science and Technology for Ceramics, 48018 Faenza, Italy  Received 29 March 2013; received in revised form 21 May 2013; accepted 23 May 2013  Available online 13 July 2013  Abstract  Ultra-high  temperature ceramics having melting points above 3500 K and high  thermal conductivities are envisaged as future receivers of concentrating solar power plants. The high pressure and solar  temperature  reactor  implemented at the focus of the Odeillo 5 kW solar furnace was used to investigate the oxidation of three refractory carbides containing different sintering additives (HfC/MoSi2 , ZrC/MoSi2 , ZrC/TaSi2 ) that could be considered as promising candidates. The concentration of the additive, TaSi2 or MoSi2 , was 20 vol%. Each kind of sample was oxidized  in air for 20 min at 1800, 2000 and 2200 K. Experiments were filmed using a video camera and the gaseous phases were analyzed in situ by mass spectrometry. Various post-test characterizations have shown that the nature of the carbide and additive strongly affects the composition of the oxide layer and therefore the high-temperature behaviour. © 2013 Elsevier Ltd. All rights reserved.  (Réacteur Hautes Pression et Température Solaire, REHPTS)  Keywords: Ceramic; SEM; XRD; High temperature corrosion; Internal oxidation  1.   Introduction  The efficiency of a concentration solar power plant highly relies on the high temperature behaviour of its solar receiver. Up to now, silicon carbide (SiC) was the only one ceramic material used  to produce various geometries for solar absorbers,1,2 but degradation of  this material becomes  relevant above 1700 K, due  to bubbles  formation and production of gaseous SiO and CO, leading to a severe mass loss of the material. Therefore SiC receivers cannot be heated at temperatures higher than 1700 K. Consequently, as for metallic receivers, an extra source of fossil energy (or biomass) has to be added after the receiver in order to end up the heating of pressurized air up to more than 1300 K and therefore insure an efficiency of energy conversion economically advantageous.3 This paper deals with  the oxidation behaviour of new ultrahigh temperature ceramics (UHTCs) keeping good mechanical properties above 2000 K,  in order  to  identify which would be  ∗  Corresponding author at: PROMES-CNRS, 7   rue du Four Solaire, 66120  Font-Romeu Odeillo, France. Tel.: +33 4 68 30 77 41; fax: +33 4 68 30 77 99.  E-mail address: ludovic.charpentier@promes.cnrs.fr (L. Charpentier).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.05.022  the best  candidates  to  elaborate new high  temperature  solar receivers. Among potential materials for such application, zirconium carbide, ZrC, presents a high melting point  (3500 K) and  interesting mechanical properties, especially  its hardness around 27 GPa, making of it one of the hardest materials among UHTC.4 On the other side, hafnium carbide, HfC, is one of the most refractory compounds available with melting point above 4100 K5-7 and  it also presents  intrinsically  spectral  selective properties.8-10 Both Hfand Zr-carbides can also be considered for  thermoionic/thermoelectric converters at high  temperature, by tuning of the grain boundary phases or carrier concentration and mobility.5-7 In spite of  their excellent properties, carbides have been hardly developed on  industrial scale due  to  the high cost of  the  raw materials and of processing and sintering.  In addition,  the main  limitation for high  temperature applications concerns the oxidation behaviour: at temperature above 1200 K the carbides start to oxidize into non protective and porous scale of ZrO2 or HfO2 according to a linear kinetics.11 Incorporation of silicon-carrying species, like SiC or transition metal silicides, was found to enable the formation of silica (ZrO2 ·SiO2 ), which was or mixed oxide  layer such as zircon  shown to improve the oxidation resistance.12 Sarin et al.13 also recently  reported  that  the  thickness of  the oxide  layer  formed        \\x0c', '2868   Table 1  L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Composition, sintering parameters, density and phases of the as-sintered ceramics tested in the REHPTS. PLS: pressureless sintering; HP: hot pressing.  Label   Initial composition  Sintering type, temperature, time,  (vol%)  environment, pressure  Final density (g/cm3 )  Final relative  density (%)  Final main phases   Composition minor  HCM   ZCM   ZCT   HfC + 20 MoSi2 ZrC + 20 MoSi2 ZrC + 20 TaSi2  PLS 2200 K, 60 min, Ar, -  PLS 2200 K, 60 min, Ar, -  HP 1970 K, 6 min, vacuum, 30 MPa   11.1  6.2   7.1   97   95   99   HfC, MoSi2 ZrC, MoSi2 ZrC, (Zr, Ta)C   phases (<5%)  (Hf, Mo)x Siy SiC, (Mo,Zr)x Siy SiC, (Ta, Zr)x Siy  on ZrB2 /SiC composite materials during oxidation in air around 1500 K decreased when the amount of SiC in the initial material increased from 0 up to 50 vol%. This study is focused on two carbides sintered with addition of Moor Ta-disilicide developed at ISTEC-CNR by pressureless sintering or hot pressing.14-16 Preliminary studies on the oxidation behaviour at high  temperature of ZrC/MoSi2 materials  in the same solar furnace facility17 gave encouraging results, thus this composite was studied in more details and the investigation was also extended to other two UHTCs, namely HfC/MoSi2 and ZrC/TaSi2 ,  in order  to understand  the effect of changing either the carbide matrix or the sintering additive.  2. Materials and methods  Materials were processed by  ISTEC-CNR  in Faenza,  Italy and  a  summary of processing parameters  and properties  is reported  in Table 1. The following starting compositions were used:  HfC + 20 vol% MoSi2 , labelled as HCM; ZrC + 20 vol% MoSi2 , labelled as ZCM; ZrC + 20 vol% TaSi2 , labelled as ZCT.  The composites containing MoSi2 , HCM and ZCM, were pressureless sintered at 2200 K  for 60 min  in flowing Ar,15,16 while the one containing TaSi2 , ZCT, was hot pressed at 1970 K for  6 min  under  30 MPa  loading. After  sintering,  the main constituent phases are  the starting ones and minor amounts of SiC  (5 vol%)  in ZrC-composites and mixed silicides, such as  Hf-Mo-Si,  in HfC. As  it  is shown  in Table 1, all  the ceramics were  almost  fully  dense with  relative  density  above  95%. In ZCT,  the  formation of  (Zr, Ta) C  solid  solutions  and of mixed  (Zr, Ta)  silicides  in  the final microstructure confirms the  cations  exchange  and  the  dissolution-reprecipitation in presence of TaSi2 .18 Therefore,  phenomena occurring  the final volumetric composition estimated through image analysis approximately resulted as:  80 HfC + 15 MoSi2 + 5 (Hf, Mo)xSiy for HCM, 80 ZrC + 13 MoSi2 + 5 (Zr, Mo)xSiy + 2 SiC for ZCM, 85 (Zr, Ta) C + 10 (Zr, Ta)xSiy + 5 SiC for ZCT.  For all the three composites, the formation of mixed silicides is a consequence of the cations mobility occurring in the liquid phase at the sintering temperatures. The formation of SiC in ZrCbased ceramics is attributed to the reaction between silicides and residual carbon present as an impurity in the starting ZrC powder (1.5 wt%). The samples were cut into 25 mm diameter and 2 mm thickness by Electrical Discharge Machining (EDM). The  reactor used  to perform high  temperature oxidation   is  the REHPTS (Réacteur Hautes Température et Pression Solaire,  High Temperature and Pressure Solar Reactor), implemented at the  focus of  the Odeillo 5 kW  solar  furnace and  it  shown  in Fig. 1. A flat mirror (heliostat), whose position  is servo-controlled to the apparent movement of the sun, reflects the incident solar flux  to a concentrator with  faceted mirrors. A shutter enables to control  the fraction of  the concentrated solar flux delivered to the sample placed inside the reactor and therefore its surface  Fig. 1.   Image of the REHPTS reactor implemented at the focus of the 5 kW solar furnace.  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2869  ±  ±  in situ   temperature. In this set-up, the sample is placed 25 mm above the focus of the solar furnace, so that very high temperatures may be obtained on a homogeneous 10 mm diameter area in the middle of the sample, with low surface gradient as UHTCs present good −1 ). thermal conductivities, and at very fast rate (up  to 100 K s Two mirrors enable a monochromatic (5  \\u242em) optical pyrometer (Ircon, Modline Plus) to measure the surface temperature in the centre zone  (6 mm) of  the sample  through a fluorine window. The pyrometer, together with all the parts present on the optical path - window and mirrors - was calibrated on a blackbody. The accuracy of  the  temperature measurement  is going  from 1400   15 K to 2100   22 K. The oxidations were performed  in air with an atmosphere continuously renewed. Due  to  the altitude (1500 m) of  the  laboratory, the total atmospheric pressure is around 87 kPa and the oxygen partial pressure pO2 is 17 kPa. The  temperature of  the samples was maintained at a constant plateau during 20 min and a video camera was used  to  follow  the oxidation process. The interest of solar concentrated energy is that the sample can be heated from ambient to the desired temperature in a few seconds. So we can clearly analyze  the very beginning of  the oxidation at one  temperature without any disturbance due  to a too long transient step. But for full information about the mechanisms of such a complicated oxidation process, a single stop of the process after 20 min could be insufficient. A mass spectrometer (Pfeiffer Omnistar) enables in situ gas phase analysis. As explained further in the text, CO is expected to be one of the main gaseous products during oxidation, but its molar weight is the same as N2 one (m/e = 28), so it is impossible to separate the contribution of CO from the one of preponderant N2 . We  therefore mainly followed  the signal corresponding  to m/e = 44, corresponding both to CO2 and gaseous SiO. The  samples were weighted before and after oxidation  in order to assess the mass variation (converted to a mass variation −2 min −1 ) and  rate expressed  in mg cm the surfaces and crosssections were analyzed after oxidation using XRD, SEM and EDS. In order to compare the stable products formed by oxidation of  the ceramics  in air,  thermodynamic calculations were performed using the GEMINI19 software. This software is based on Gibbs energy minimization in order to define which compounds are  the most stable at  thermodynamic equilibrium, according to the initial conditions (temperature, pressure, molar fractions of reactants). The COACH database management system provides to GEMINI the requested thermodynamic data, built from the JANAF database,20 for each possible compound. Chemical kinetics and nucleation energies are not  taken  into account by GEMINI software.  3. Results  3.1. Thermodynamic calculations  Fig. 2 presents  the moles of  solid phases  and  the molar fraction of gaseous products CO, CO2 and SiO and molybdenum oxides,  from 1000  to 2200 K  for  the  three carbides as a  function of temperature at the thermodynamic equilibrium with the following initial conditions:   10 moles of HCM/ZCM/ZCT,  150 moles of air (120 moles N2 , 30 moles O2 ), Total pressure: 105 Pa.  The molar compositions of HCM, ZCM and ZCT were fixed according to the compositions of the as-sintered materials given in Section 2 and the available thermodynamic data:  HCM: 86.5 HfC and 13.5 MoSi2 , mol%; ZCM: 85.6 ZrC, 9 MoSi2 , 2.8 ZrSi2 and 2.6 SiC, mol%; ZCT: 65 ZrC, 24.1 TaC, 5 ZrSi2 and 4.9 SiC, mol%.  ϕ  −  For simplicity, the (Zr, Ta) C solid solution of ZCT has been decomposed  into ZrC and TaC phases, which allows a  rough calculation. The number of elements N (Zr or Hf, C, Ta or Mo, Si, N, O) is 6 in any of the calculations. As there cannot be more than   = 4 phases (3 solid oxides: ZrO2 or HfO2 , MoO3 or Ta2O5 , silica or zircon + 1 gaseous phase) coexisting at any temperature, the variance (v = N + 2  ϕ) is 4. The fixed three independent parameters are total pressure, initial composition of the gaseous phase (air) and  initial composition of each solid material, hence only one  thermodynamic equilibrium  is possible for each  temperature. We have chosen to use 150 moles of air in order to provide enough oxygen (i.e. 30 moles) to oxidize completely 10 moles of material and therefore to reach the thermodynamic equilibrium for all of  the possible  reactions. These proportions may  look huge.  Indeed, oxidizing 1 mole of material under 15 moles of air - or lower proportion respecting the 1:15 ratio - would give the same conclusions, we have chosen so high proportions  in order to have more significant amounts of products and clearer interpretations, neglecting fewer minor products. Due to the oxidation kinetics, we do not expect  to reach  the  thermodynamic equilibrium during experiments  that will  last only 20 min. But the thermodynamic calculations give an indication on what the final state of  the system should be, enabling us  to estimate  the stability of one potential material in extreme conditions. Looking at  the solid species formed  in  the composites containing MoSi2 ,  in Fig. 2a and c, HCM should be oxidized  into hafnia and silica at any  temperature. ZCM should form zirconia and zircon up  to 2000 K, and above  this  temperature only zirconia and silica should be stable. Molybdenum solid oxide, MoO3 , is stable only at 1000 K. Moving to the volatile species in HCM and ZCM,  the plots  in Fig. 2b and d have analogous trends:  the main gaseous product  is CO2 whose molar fraction remains between 6% and 7%,  then O2 in  the atmosphere  in both cases, showing there was enough oxygen to complete any potential oxidation. Other gaseous products are CO, SiO, and molybdenum oxides. These last have a general formula MonO3n with n = 1-5. The  increase  in  the molar fraction of MonO3n is due  to  the  fact  that higher  temperatures  favour  the  formation of “lighter” oxides, Mo3O9 being preponderant  from 1100  to 1800 K, MoO3 from 2000 K. The molar  fraction of molybdenum oxides  is a bit  lower  for ZCM  than  for HCM, which  is    \\x0c', '2870   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Fig. 2. Amount of solid phases (left) and molar fraction of   the main gaseous products (right) at   the   thermodynamic equilibrium as a function of   the   temperature,  calculated using GEMINI software: (a, b) for HCM, (c, d) for ZCM and (e, f) for ZCT.  expected as HCM  is  richer  in MoSi2 than ZCM. The  interest of using ZCT  instead of ZCM or HCM could be  the possibility  to form  tantalum solid oxide Ta2O5 at any  temperature, as Fig. 2e  shows. Consequently, besides CO2 , CO and SiO, no other gaseous oxide would evolve  in  the  investigated  temperature range (Fig. 2f). The oxygen consumption is a bit lower, due to the fact that the formation of Ta2O5 requires fewer moles of oxidant  than  the  formation of MonO3n .  In all  the  three cases, one can see that conspicuous CO escape occurs around 2000 K. From these thermodynamic calculations, it seems that ZCT is the most promising composite, due  to  the formation of more solid products and  less gaseous species  than HCM and ZCM. Since the  formation of solid phases has slower kinetic as compared to gases, ZCT is expected to undergo less degradation than the other two ceramics.  3.2.   In situ analyses  3.2.1. Video images  Fig. 3 presents video captions of  the various materials after 15 min oxidation tests in air at 1800, 2000 and 2200 K, respectively. We can observe that at 1800 K, the oxide layer that forms  ×  ×  vs. 5.8   ×  ×  −1  −1 vs. 8   on HfC is rough and broken into several fragments, whereas the oxide layer that forms on ZrC materials looks quite smooth and well adherent  to  the carbide. This could be explained by a difference  in  the  thermal expansion coefficient: HfC has a higher linear coefficient of thermal expansion than HfO2 (up to 1670 K: −6 K −6 K −1 ,  7.3   10  10 respectively),21,22 so  the oxide  layer  tends  to break due  to  the expansion of  the carbide. On  the other hand, ZrC has a  lower coefficient  than ZrO2 (up −6 K −6 K −1 , respectively),21,22 to 1670 K: 7.6   10  10 that means that the oxide could dilate more than the carbide and results in a complete coverage of the carbide. At 2000 K, HCM still appears rough and porous and has a bigger volume than the starting disc. We can also notice  that HCM has slightly moved from the support due to the release of gaseous products also on the back face. ZCM clearly shows the formation of bubbles on the top, probably due to the formation of liquid silica, and ZCT is the least damaged and has an aspect not notably different from the photo taken at 1800 K. At 2200 K, the composites containing MoSi2 , HCM and ZCM, display clear boiling phenomena, while for ZCT just a little bump can be observed on the upper surface. Therefore, according  to  these visual  inspections,  it seems  that ZrC has better oxidation resistance than HfC, and that TaSi2 has  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2871  Fig. 3. Video captures of the three UHTC after 15 min. Oxidation in air at (a) 1800 K, (b) 2000 K and (c) 2200 K.  more beneficial effect  than MoSi2 ,  in agreement with  the  thermodynamic calculations that showed that the oxidation of ZCT produces more solid oxides and less gaseous products than the ZCM and HCM.  3.2.2. Mass spectrometry  Fig. 4 displays  the evolution of CO2 and/or SiO concentration as a  function of  the oxidation  temperature  for  the  three carbides, determined using mass spectrometry. We can observe that, at 1800 and 2000 K, HCM is the material that releases more gaseous products during oxidation, in agreement with the video images (Fig. 3). This can be due to the fact that its oxide scale is subjected to continuous rupture and allows the gases escape from the oxide/carbide interface. Nevertheless, after about 10 min at 2000 K and from  this  temperature (Fig. 4b and c),  the slope of the HCM curve is less steep than at the beginning of the oxidation. The apparition of a liquid glass and its vigorous migration to  the surface  in droplet shape, as confirmed  in Fig. 3, could partially fill  the porous structure and prevent gaseous products escape. Equivalent amount of CO2 and SiO gases is detected in ZCM and ZCT during  the oxidation at 1800 K and during  the first 10 min of oxidation at 2000 K. During  the second half of the oxidation at 2000 K, ZCT evolves higher amount of gaseous products. This difference can be due to the fact that during sintering, (Zr, Ta) Si2 mixed silicides are formed and they are less stable than MoSi2 at high temperature. ZrSi2 has a melting point around 1900 K23 and has  the  tendency  to decompose around  these  temperatures, releasing more SiO. A second explanation comes  from  the video observations:  the  liquid oxide migrates to the surface of ZCM and can therefore prevent gaseous products  from being  released. Reduced amount of silicides  in  the as-sintered ZCT seems  to  limit  the  formation of big bubbles inside  the  liquid oxide,  therefore gaseous products may faster escape  through fissures and  the bursting of  smaller bubbles. Both hypotheses could also explain why at 2200 K  the gases escaping  trend changes: ZCT  releases more CO2 and/or SiO than  the  two other materials, which have now a comparable behaviour. We observe  that  the higher  is  the  temperature,  the higher is the amount of gaseous products released by ZrC-based materials.  3.3. Post-experimental characterizations  3.3.1. Mass variation rates  Table 2 gives  the global mass variation rate for each mate−2 min −1 . In general,  rial expressed  in mg cm the higher  is  the temperature,  the faster  is  the oxide growth and  the mass gain. The mass variation rate of HCM could not be evaluated owing to  important breaking and  removal of  the oxide  layer on HfC during  the experiments. The nature of  the additive  in  the ZrC materials has  little  influence at 1800 K, but at higher  temperature,  the mass variation  rate  is more  important when TaSi2 is used. Tantalum seems therefore to play a significant role in the  Table 2  Average mass variation rates and crystalline phases obtained by XRD for each material as a function of the oxidation temperature.  Mass variation rate −2 min −1 ) (mg cm  Crystalline phases by  XRD  Temperature (K)   1800   2000   2200   1800   2000   2200   HCMa  n.a.   n.a.   n.a.   ZCM   1.28   1.72   2.23   ZCT  1.20  2.37  3.86  HfO2 , HfSiO4 traces  HfO2 , HfSiO4 HfO2 , HfSiO4  ZrO2 , MoSi2 , Mo5 Si3 traces  ZrO2 ZrO2  TaZr2.75O8 , ZrO2 TaZr2.75O8 , ZrO2 TaZr2.75O8 , ZrO2  a Loss of external surface during test.  \\x0c', '2872   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  0  0.3  0.6  0.9  1.2  1.5  0  20 0  40 0  60 0  80 0  100 0  12 00  Time (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  0  0.3  0.6  0.9  1.2  1.5  0  200  400  600  800  1000  1200  T ime (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  0  0.3  0.6  0.9  1.2  1.5  0  20  0  40  0  60  0  80  0  100  0  120  0  T ime   (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  a)  b)  c)  HCM  ZCM  ZCT  Fig. 4. Measured concentration of m/e = 44   (SiO and CO2 produced) at  during  the  oxidation  of  the UHTCs  (a)  1800 K,   (b)   2000 K   and   (c)   2200 K   specimens.  0   500   1000   1500   2000   2500   3000   3500   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  2θ  (°)   2θ  (°)   2θ  (°)   t-HfSiO4   m-HfO  2   a)   0   500   1000   1500   2000   2500   3000   3500   4000   4500   5000   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  m-ZrO2   b)   0   500   1000   1500   2000   2500   3000   3500   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  m-ZrO2   o-TaZr2.75O8   c)   Fig. 5. XRD patterns of   the samples oxidized during 20 min   in air at 2000 K:  (a) HCM, (b) ZCM and (c) ZCT.  oxide  formation, especially at 2200 K, as SEM  will demonstrate.  investigations  3.3.2. X-ray diffraction  X-ray diffraction results performed on the surface of the three composites after oxidation  tests are  summarized  in Table 2. Fig. 5 compares the diffraction patterns for the samples oxidized at 2000 K, for each kind of material. The reference patterns used for the indexation of the peaks are: 83-0944 (monoclinic ZrO2 ), 65-1142 (monoclinic HfO2 ), 77-1759 (tetragonal HfSiO4 ) and 42-0060 (orthorhombic TaZr2.75O8 ). It has to be underlined that in HCM the external layer was detached and some pieces were lost during  the  tests at all  temperatures,  so  the X-ray analyses refer  to  the remaining exposed surface (Fig. 5a). The main crystalline phases  left at 1800 K  in HCM are monoclinic hafnia and traces of tetragonal hafnium silicate (hafnon), HfSiO4 . As  the oxidation  temperature  increases,  the amount of hafnon increases as well. Since hafnon decomposition  is  reported  to                                                                                                                            \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2873  occur above 2000 K24 that  is around  the  temperatures  tested in  the present work, we can explain  increasing amount of  this phase at higher temperatures probably in relation to the removal of a thicker and thicker outermost layer. In the ZCM composite (Fig. 5b), no traces of zircon (ZrSiO4 ) were detected at any temperature and  the main crystalline phases  is monoclinic ZrO2 with an evident preferential orientation along  the  (0 0 2)  lattice plane at  the highest oxidation  temperature. MoSi2 peaks, cubic ZrO2 and traces of Mo5Si3 are visible only at 1800 K. At all  temperatures,  the main crystalline phase of ZCT  is a mixed oxide with composition TaZr2.75O8 with an orthorhombic structure with preferred orientation along  the  (0 2 0)  lattice plane (Fig. 5c); its peaks become sharper while increasing the temperature,  indicating  improved crystallization. Since  the scattering coefficient for the orthorhombic TaZr2.75O8 phase has not been published  in  the  literature or  in available  ICSD database, we considered  this phase  as  a  solid  solution between 0.5 moles orthorhombic Ta2O5 (#54-514) and 3 moles  tetragonal ZrO2 (#42-1164) and estimated a scattering coefficient of 4.7. Monoclinic ZrO2 is also present and  its peaks signal  increases with temperature up to about 25 vol% at 2200 K, indicating the higher stability of pure oxide over  the mixed one. Cubic and  tetragonal ZrO2 could be present  too, but  the superimposition of  the main peaks with the mixed oxide hinders a conclusive analysis. Other authors25,26 and similar  tests on ZrC-based composites showed  indeed  that carbon, coming  from  the oxidation of  the carbide, stabilizes c-ZrO2 at low temperatures.  3.3.3. SEM and EDS analyses  Fig. 6 presents SEM  images of  the sample surfaces  in  the central  region where  temperature was measured  for  the  three composites at 1800, 2000 and 2200 K. As a rule of thumb, dark regions correspond to silica-based glass, bright phases to heavier oxides. The  residual external  surface of HCM  looks  similar after oxidation at 1800 and 2000 K, composed by porous and cracked HfO2 , with bright contrast, and HfSiO4 , with light grey contrast and smoother appearance. At 2200 K, most of the surface is based on HfO2 and HfSiO4 , but large areas are covered with dark silica-based glass  (Fig. 6a). These observations are in agreement with the video images, revealing the vigorous formation of bubbles at 2200 K, and with  the X-ray diffraction. For ZCM (Fig. 6b), at 1800 K,  the surface  is made of cracked ZrO2 and some MoSi2 phase can be observed, with an irregular shape and smooth edges. Increasing  the oxidation  temperature to 2000 K, silica migration to the surface occurs. ZrO2 phase is present in two morphologies: as cracked large grains (10  \\u242em) and brain-like smaller grains,  indicating precipitation, nucleation and growth  from  liquid phase. At  this  temperature,  large areas of the sample surface were covered with dark silica-based glass, accordingly to video images. Further temperature increase induced the formation of larger flat ZrO2 grains, deriving from the melting and recrystallization of  this phase  that  left silica at the grain boundaries. The addition of TaSi2 to a ZrC matrix generated a variety of articulate and impressive morphologies at the various oxidation temperatures (Fig. 6c). At 1800 K the surface has a rough cracked aspect with TaZr2.75O8 and ZrO2 being the  \\u242em   main phases. The surface of ZCT at 2000 K is mainly composed of petal-like grains of TaZr2.75O8 which  form volcanoes and tend to microfissuration. At 2200 K melting and recrystallization of  the mixed oxide occurred and 20  large grains precipitated  leaving  residual silica, containing Zr  traces, at  the grain boundaries. These grains are  in  turn composed of polyhedral structures of ZrO2 and TaZr2.75O8 showing  the growth planes elegantly decorated by a dendritic irregularly shaped phase identified as ZrO2 and ZrSiO4 containing small traces of Ta. In the ZrC-based composites, besides  these  tiny shapes  in ZCT oxidized at 2200 K, no  further ZrSiO4 has been clearly detected either by XRD or SEM-EDS. As a matter of  fact,  the  formation of zircon from mixed oxides starts at around 1470 K, with highest formation rate between 1770 and 1820 K.27 At higher temperatures ZrSiO4 concentration decreases again  indicating the occurrence of the decomposition reaction of zircon, reported to start at 1820 K, which accelerates with  increasing  temperatures. Considering the quenching of the samples at the end of the test, we can reasonably understand that upon cooling, the reverse reaction from ZrO2 and SiO2 to form zircon is suppressed. Fig. 7 presents  the cross-sections of HCM, ZCM and ZCT oxidized at 2000 or 2200 K. Fig. 7a-c shows  the morphology of  the  remaining HCM oxidized at 2000 K. Keeping  in mind that HCM lost the upper part, presumably composed by porous HfO2 , the remaining surface scale is composed by dense HfO2 , run through by HfSiO4 and MoSi2 oxidation products, that are Mo5Si3 phase and SiO2 glass (Fig. 7b). Moving inward (Fig. 7c), we find a mixed oxy-carbide and MoSi2 with its oxidation products.  It has  to be noticed  that, despite  the external  layer was detached, these remaining ones are crack free and well adherent to the unoxidized bulk. Given the abundant presence of hafnon, we can deduce  that  the  left surface did not exceed  the decomposition  temperature of  this compound,  reported  to be above 2000 K.24 Fig. 7d-g shows the morphology of ZCM oxidized at 2000 K. An external amorphous SiO2 phase, not homogeneously distributed  (as  result of dropleting phenomena), covers ZrO2 grains which have brain-like shape and  together  they  form a compact dense layer (Fig. 7e). Moving inward, a porous and brittle ZrO2 layer separates a dense region composed by ZrC grains surrounded by granulous Zr-C-O, MoSi2 and Mo5Si3 with SiO2 thin glassy film and droplets (Fig. 7f). As one approaches  the bulk (Fig. 7g), SiC particles are found  instead of SiO2 and  the oxidation mechanisms of matrix and MoSi2 are evident: MoSi2 is surrounded by brighter Mo5Si3 , adjacent  to  the pre-existing SiC. These islands are in turn contoured by Zr oxy-carbide which encompasses  the original ZrC grains. These observations  for ZCM and HCM are in agreement with the mechanism proposed by Shimada28 : an  intermediate oxy-carbide forms close  to  the interface with  the carbide, until  it  is completely evolved  to  the oxide phase. ZCT oxidized at 2000 K  (not shown) has a  layered structure with the outermost scale composed of TaZr2.75O8 with some  inclusions of Si-C-O phases.  In  the second  layer, TaZr2.75O8 is still  the main phase, but as secondary phase we find already SiC instead of oxides. Moving forward towards the bulk, at around 400  \\u242em from the surface, a complex mixture of (Zr, Ta)Si2 , (Zr, Ta)-oxy-carbide and SiC phases are found. The complex  (Zr, Ta)-C-O phase  is  the  result of partial oxidation  \\x0c', '2874   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Fig. 6. SEM images of the surface of (a) HCM, (b) ZCM and (c) ZCT after 20 min. Oxidation in air at 1800, 2000 and 2200 K.  of  the starting matrix grain, composed also by (Zr, Ta) C (see Table 1). For ZCT composite, it is interesting to thoroughly analyze  the morphology of  the oxidized scale after  test at 2200 K, which marks the limit temperature between the goodness of ZCT as compared  to ZCM. Fig. 7h-m  shows  the cross-section of ZCT oxidized at 2200 K, where all  the  thickness of  the sample  resulted modified. The outermost  thick scale  is composed by a compact ZrO2 layer where 20-30  \\u242em large porosities can be  found  (Fig. 7i).  In  this  region, SiO2 droplets  (dark) containing Zr  traces are surrounded by a brighter phase,  identified as a solid solution with possible formula (Zr0.8Ta0.2 )O2 . Right beneath this layer, the mixed TaZr2.75O8 oxide is present forming a  thick dense scale  including  isolated porosity (Fig. 7l). Moving further  inward, fine grained ZrO2 with Ta  traces and SiO2 discrete phases stand above the already mentioned (Zr, Ta)-oxycarbide and SiC phases (Fig. 7m, upper part). The centre of the sample is composed of a mixed (Zr, Ta)-oxy-carbide containing lower amounts of oxygen (Fig. 7m, lower part). This articulate morphology is a sign of complex oxidation mechanisms, including melting, phase separation and reprecipitation, occurring at 2200 K.  4. Discussion  In  light of  the multiple analyses performed during and after oxidation  in  the  temperature  range 1800-2200 K on Hfand ZrC-based materials, we can deduce several observations.  4.1. Background on the oxidation behaviour  Numerous reports have been devoted to the study of the oxidation of carbides of groups  IV-VI and  the most widespread outcomes is that transition metal carbides, generally referred to as MeC, oxidize according to the following reactions:  O2 →  MeCxOy +  C1−x + MeC  MeCxOy +  C1−x +  O2 →  MeO2 +   O2−y   COy  +  (1)  (2)  These general reactions can be extended to the case of ZCT, where a solid solution  is oxidized  to  the corresponding mixed oxy-carbide and then to the oxide (3):  O2 →  Ta)CxOy →   TaZr2.75O8   Ta)C   (Zr,  (Zr,  +  (3)    \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2875  Fig. 7. SEM   images of   the polished cross-section of   (a-c)   for HCM,   (d-g)   for ZCM after 20 min. Oxidation   in air at 2000 K and   (h-m)   for ZCT after 20 min.  Oxidation in air at 2200 K.  in,25,26,28-30  According  to  the  studies  reported  oxide growth of carbides  is parabolic  indicating protective oxidation behaviour up  to 2000 K. During  the oxidation, gaseous CO  is produced as a by-product. CO generation  introduces porosity which allows diffusion via pores. Gas formed below  the oxide layer can also lift and disrupt the oxide layer, like in the case of HCM. Nevertheless,  the generated oxide  layer  is partially protective, as testified by the parabolic rate of gases evolution which is believed to control the oxide growth. A different behaviour is observed upon oxidation at higher  temperatures, as  illustrated by previous studies.30,31 At higher  temperatures,  the evolution of gases becomes less important as the oxide layer starts to sinter significantly and  the effective volume available for gaseous diffusion  is  reduced. This  implies  that  the sintering  temperatures of the surface oxides determine a change in the oxidation behaviour. The addition of significant amounts of silicides can partially alter  the oxidation behaviour, so  the oxidation  reactions involving the silicides and the other secondary phases, like SiC, should be considered as well. Reactions  (4),  (6) and  (8) are referred to as “passive oxidation reactions”, with formation of protective SiO2 glass, while  reactions  (5),  (7) and  (9) are  generally defined as “active oxidation reactions”, owing  evolution of gaseous species (SiO and CO): 5MoSi2 +  7O2 →  Mo5Si3 +  7SiO2 MoSi2 +  (5/2)O2 →  MoO3 +  2SiO  2TaSi2 +  (13/2)O2 →  Ta2O5 + 2TaSi2 +  (9/2)O2 →  Ta2O5 +  2O2 →  SiO2 +  O2 →   4SiO2   4SiO   +  +  SiC   SiC    CO2  +   SiO    CO   to   the  (4)  (5)  (6)  (7)  (8)  (9)  These general reactions can be extended to the case of mixed silicides. Passive oxidation  is expected  to occur at  temperature up to about 1800 K, while active reactions occur at higher temperatures and in low oxygen partial pressure. The surface of the samples in the present tests falls in the region of active regimes, that  is  from 1800 K on. Solid  reaction products are observed in the samples cross-sections, owing to different oxygen partial pressure and temperature gradient. In principle, oxidation of silicides should produce silica that diffuses through the surface and  \\x0c', '2876   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  forms a stable and continuous silica  layer. However, previous studies on  the oxidation of  these composites have shown  that even at temperatures lower than 1500 K, no protective oxidation layer was observed on the surface.15,16 This can be due to several reasons:  there  is not enough silica  to fill all  the volume expansion associated with the formation of the porous oxidation layer and/or large CO escape resulting from oxidation of the carbide can favour the dissociation of silica to gaseous SiO. Although no continuous silica layer was found on the surface, the presence of partially filled porosity  in  the cross-section can hinder  the fast diffusion of gaseous species towards the unreacted bulk.  4.2. Effect of the carbide matrix  First, we could say  that  thermodynamic calculations are  in quite good agreement with post-test analyses, confirming  the formation of the main solid products foreseen for the tabulated compounds. Here we will not consider  the formation and stability of HfSiO4 , owing  to  the uncertainty of  the real external layer. In addition,  it has  to be underlined  that  thermodynamic data of the mixed oxides formed in ZCT, TaZr2.75O8 , are not well known and  therefore could not be  included  in  the calculations. Moreover, the calculations considered only the thermodynamic equilibrium and did not  take  into consideration  the kinetic of the reactions or the energy required for creating new interfaces, which  is  the  reason why  the  solid compounds observed are slightly different  from  those predicted by  the  thermodynamic calculations. Nevertheless,  the outermost oxide  layers  involve Hf and Si elements (HCM), Zr and Si elements (ZCM), Zr, Si and Ta elements (ZCT), which is coherent with the thermodynamic results. As expected  from  these calculations, no molybdenum solid oxides appeared, but Mo5Si3 phase was  found close  to the interface with the carbide inside HCM and ZCM materials, probably resulting from incomplete oxidation of the additive. Video images evidenced quite surprising phenomena, as HfCbased ceramic resulted notably worse  than ZrC-based ones, as reported by other authors.25 Mass spectroscopy analyses also confirmed  the consistent oxidation of HCM over ZrC-based composites up  to 2000 K  involving great amount of gaseous species. However, at 2200 K,  this picture changed and HCM showed a behaviour  similar  to ZCM, probably owing  to  the achievement of equilibrium between gas production and glass boiling on the top. Of course these are just preliminary results of very complex mechanisms  that need  to be further  investigated and thoroughly confirmed.  4.3. Effect of the sintering additive  The addition of TaSi2 rather  than MoSi2 has  further consequences on  the oxidation behaviour of ZrC.  In addition,  it is  important  to underline  that  the  total amount of silicides  in the final microstructure was notably reduced from  the starting 20 vol%. Although TaSi2 has a high melting point  (2470 K), during sintering and oxidation mixed silicides (Zr, Ta)Si2 and Si-based phases with  significantly  lower melting point were formed. On  the other hand, upon oxidation,  a mixed oxide with stoichiometry TaZr2.75O8 was formed. Some authors have  studied  the effect of Ta-addition on  the oxidation behaviour of ZrB2 , where the introduction of tantalum was for increasing the glass viscosity and thus limiting the oxygen diffusivity.32 However, some authors also recognized that the melting point of this phase could be significantly  lower  than  that of zirconia33 with consequences on the high temperature stability of the oxide. This hypothesis seems to find a confirmation in the test performed on ZCT at 2200 K, where evident melting, dissociation, evaporation and re-precipitation phenomena occurred  leaving mainly ZrO2 in  the external  layer of  the final microstructure. According  to Bhattacharya et al.,34 who computed a recent phase diagram for the ZrO2-Ta2O5 system, Zr-Ta-O solution melts from 2100 K, that could explain why no deformation is observed up to 2000 K. The evident consequence is the better aspect of the ZCT samples as  compared  to  the ZCM ones  (Fig. 3),  suggesting  a more beneficial effect of  the addition of TaSi2 , compared  to MoSi2 . It  is probable  that  the petal-like open structure of ZCT and  the complex architecture of the modified layer allow for controlled gaseous  products  release,  avoiding  excessive  pressurization inside  the sample,  thus preserving  the scale from spalling and delamination. Only at 2200 K, ZCT showed some deformation owing  to  the  formation of  a dense  thick ZrO2 layer, which favoured the formation of large gas pockets below this scale.  4.4. Oxidation mechanisms in the different temperature ranges  According to the analyses performed and the studies carried out, we can now delineate the most important phenomena occurring during oxidation of these carbides in the temperature range between 1800 and 2200 K.  1800 K: At  this  relatively  low  temperature,  the  oxidation behaviour of ZrC-based composites is better than HfC, due to that fact that the ZrO2 -based oxide formed has the tendency to shrink and the effective volume available  for gaseous diffusion  is  reduced. On  the other hand, at this temperature, the oxidation of HfC is still dominated by  the escape of gases  through  the highly porous oxide layer. Gas evolution data suggest a linear rather than parabolic behaviour for all the composites. 2000 K: The oxidation of HfC/MoSi2 is still dominated by the escape of gases through the highly porous oxide layer. For the ZrC composite containing MoSi2 , data in Fig. 4 suggest a parabolic behaviour. In this case, it is likely that sintering of the ZrO2 oxide becomes important and protects the composite from diffusion of oxygen. The composite containing TaSi2 has a linear behaviour up to 600 s. After that time, there is an abrupt increase of gas production, that is probably related to the evolution of the cited low melting phases, especially (Zr, Ta)xSiy . At this temperature, sintering occurred for zirconia scale but not yet for the hafnia oxide layer. 2200 K: At this temperature, the HfC/MoSi2 composite shows a marked improved behaviour in comparison with the lower  temperatures, which could be attributed  to  the sintering of  the hafnia  scale. Data on gas evolution  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2877  indicate  that  the extent of oxidation activity  is much reduced at 2200 K. Moreover, the parabolic behaviour recorded for  the gas evolution  is comparable or even better  than  that of ZrC/MoSi2 composite. In contrast, the ZrC/TaSi2 composite changes  its behaviour from protective  to non-protective. This  is due  to  the  instability of  the external TaZr2.75O8 phase at 2200 K and indicates that this oxide is not so effective in hindering the oxygen diffusivity. This results in enhanced incorporation and diffusion of oxygen down  to  the core of the sample, as well as release of gaseous species, with a linear increase with time.  5. Conclusions  Oxidation studies on Hfand Zr-carbide composites sintered with MoSi2 or TaSi2 were performed  in  the high  temperature range of 1800-2200 K  in order  to discriminate  the oxidation performances of the composites depending on matrix, sintering additive and temperature regime. Oxidation of HfC/20 vol% MoSi2 occurs with  formation of  gaseous  products  and  a  fragile  and  brittle HfO2 layer, accompanied by a notable volume expansion which causes  its delamination from  the bulk already at 1800 K. At 2200 K conspicuous melting of silica, deriving  from  the oxidation of  the silicide, and  its migration  to  the surface  take place with  formation of glass drops on the sintered HfO2 external scale. The same silica drops were observed  in ZrC/20 vol% MoSi2 composite from 2000 K on. At higher temperatures zirconia melting occurred with consequent limitation of gaseous products escape. Changing  the additive of  the ZrC matrix and using 20 vol% TaSi2 appears  to reduce  the deformation up  to 2000 K,  thanks to  the formation of a mixed oxide (TaZr2.75O8 ) partially filled with silica. Its peculiar open solid structure immersed in viscous glass allows the system to accommodate the new oxide structure without notable  sample distortion and, at  the  same  time,  the continuous glass hinders gases escape or oxygen penetration. Problems occur when the temperature achieves 2200 K: melting of the mixed oxide occurs, vigorous evaporation of CO and SiO is allowed and oxygen is let flowing all across the sample. The oxidation mechanisms of  these carbide composites are a very complex matter  indeed and  this study does not claim  to be conclusive, however it pointed out that there is not a material suitable for all  the  temperature conditions: ZrC + TaSi2 seems more adequate  for  temperatures up  to 2000 K, HfC + MoSi2 has better performances at  temperatures  from 2200 K on and ZrC + MoSi2 has  intermediate behaviour between  the  two. A following step could be the design of a material containing low amount of manifolds phases: TaSi2 to induce glass immiscibility  in medium-high  temperature regime and HfC for oxidation protection in ultra-high temperature regime. ZrC-based materials could be used to design future solar receivers if the working temperatures will not overpass 2000 K. HfC does not seem adequate due  to  the  formation of brittle solid particles  that could be drained  in  the  coolant  and may  erode  the  tubes  and  the turbine.  References  1. 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The   Zr-Si   system.   Bull   Alloy   Phase Diagr   1990;11:  513-9.  \\x0c', '2878   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  24. Monaghan S, Greer   JC, Elliot SD. Thermal decomposition mechanisms  of  hafnium  and  2005;97:114911.  zirconium   silicates   at   the   atomic   scale.   J Appl Phys  30. Bargeron CB, Benson RC, Jette AN, Phillips TE. Oxidation of Hafnium        C   to 2060  C. J Am Ceram Soc  carbide  in  the  temperature  1993;76:1040-6.  range 1400  25. Voitovich F, Pugach EA. High-temperature oxidation of ZrC   and HfC.  31. Courtright EL, Prater JT, Holcomb GR, St Pierre GR, Rapp RA. Oxidation  Poroshkov Metall 1973;11:67-74.  26. Shimada   S,   Inagaki M,   Suzuki M. Microstructural   observation   of  of hafnium carbide and hafnium carbide with additions of  praseodymium. Oxid Met 1991;36:423-37.  tantalum and  formed   by   oxidation   of   ZrC.   J Mater   Res  32. Talmy IG, Kaykoski JA, Opeka MM. High temperature chemistry and oxi the  ZrC/ZrO2 interface  1996;11:2594-7.  27. Kaiser A, Lobert M, Telle R. Thermal stability of zircon. J Eur Ceram Soc 2008;28:2199-211.  28. Shimada S. Interfacial reaction on oxidation of carbides with formation of carbon. Solid State Ionics 2001;141-142:99-104.  29. Rudneva VV, Galevskii GV. Investigation of   thermal oxidation resistance  of nanopowders of  2007;48:143-7.  refractory carbides and borides. Russ J Non-Ferr Met  dation of ZrB2 ceramics containing SiC, Si3N4 , Ta5 Si3 , and TaSi2 . J Am Ceram Soc 2008;91:2250-7.  33. Levin EM, Robbins CR, McMurdie HF. Phase diagram   for   ceramists.  Columbus: The American Ceramic Soc., Inc.; 1964.  34. Bhattacharya AK,   Shklover V,   Steurer W, Witz G,   Bossmann HP,  Fabrichnaya O.   Ta2O5 -Y2O3 -ZrO2 thermodynamic  description.   system:   experimental   J   Eur Ceram   Soc   study  and 2011;31:  preliminary   249-57.  \\x0c']"
},{
  "_id": 84,
  "PDF": "High temperature oxidation of Zr- and Hf-carbides- Influence of matrix and sintering additive.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 33 (2013) 2867-2878  High temperature oxidation of Zrand Hf-carbides: Influence of matrix and sintering additive  Ludovic Charpentier a,∗  , Marianne Balat-Pichelin a , Diletta Sciti b , Laura Silvestroni b  a PROMES-CNRS, Laboratoire Procédés, Matériaux et Energie Solaire, 66120 Font-Romeu Odeillo, France b CNR-ISTEC, Institute of Science and Technology for Ceramics, 48018 Faenza, Italy  Received 29 March 2013; received in revised form 21 May 2013; accepted 23 May 2013  Available online 13 July 2013  Abstract  Ultra-high  temperature ceramics having melting points above 3500 K and high  thermal conductivities are envisaged as future receivers of concentrating solar power plants. The high pressure and solar  temperature  reactor  implemented at the focus of the Odeillo 5 kW solar furnace was used to investigate the oxidation of three refractory carbides containing different sintering additives (HfC/MoSi2 , ZrC/MoSi2 , ZrC/TaSi2 ) that could be considered as promising candidates. The concentration of the additive, TaSi2 or MoSi2 , was 20 vol%. Each kind of sample was oxidized  in air for 20 min at 1800, 2000 and 2200 K. Experiments were filmed using a video camera and the gaseous phases were analyzed in situ by mass spectrometry. Various post-test characterizations have shown that the nature of the carbide and additive strongly affects the composition of the oxide layer and therefore the high-temperature behaviour. © 2013 Elsevier Ltd. All rights reserved.  (Réacteur Hautes Pression et Température Solaire, REHPTS)  Keywords: Ceramic; SEM; XRD; High temperature corrosion; Internal oxidation  1.   Introduction  The efficiency of a concentration solar power plant highly relies on the high temperature behaviour of its solar receiver. Up to now, silicon carbide (SiC) was the only one ceramic material used  to produce various geometries for solar absorbers,1,2 but degradation of  this material becomes  relevant above 1700 K, due  to bubbles  formation and production of gaseous SiO and CO, leading to a severe mass loss of the material. Therefore SiC receivers cannot be heated at temperatures higher than 1700 K. Consequently, as for metallic receivers, an extra source of fossil energy (or biomass) has to be added after the receiver in order to end up the heating of pressurized air up to more than 1300 K and therefore insure an efficiency of energy conversion economically advantageous.3 This paper deals with  the oxidation behaviour of new ultrahigh temperature ceramics (UHTCs) keeping good mechanical properties above 2000 K,  in order  to  identify which would be  ∗  Corresponding author at: PROMES-CNRS, 7   rue du Four Solaire, 66120  Font-Romeu Odeillo, France. Tel.: +33 4 68 30 77 41; fax: +33 4 68 30 77 99.  E-mail address: ludovic.charpentier@promes.cnrs.fr (L. Charpentier).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.05.022  the best  candidates  to  elaborate new high  temperature  solar receivers. Among potential materials for such application, zirconium carbide, ZrC, presents a high melting point  (3500 K) and  interesting mechanical properties, especially  its hardness around 27 GPa, making of it one of the hardest materials among UHTC.4 On the other side, hafnium carbide, HfC, is one of the most refractory compounds available with melting point above 4100 K5-7 and  it also presents  intrinsically  spectral  selective properties.8-10 Both Hfand Zr-carbides can also be considered for  thermoionic/thermoelectric converters at high  temperature, by tuning of the grain boundary phases or carrier concentration and mobility.5-7 In spite of  their excellent properties, carbides have been hardly developed on  industrial scale due  to  the high cost of  the  raw materials and of processing and sintering.  In addition,  the main  limitation for high  temperature applications concerns the oxidation behaviour: at temperature above 1200 K the carbides start to oxidize into non protective and porous scale of ZrO2 or HfO2 according to a linear kinetics.11 Incorporation of silicon-carrying species, like SiC or transition metal silicides, was found to enable the formation of silica (ZrO2 ·SiO2 ), which was or mixed oxide  layer such as zircon  shown to improve the oxidation resistance.12 Sarin et al.13 also recently  reported  that  the  thickness of  the oxide  layer  formed        \\x0c', '2868   Table 1  L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Composition, sintering parameters, density and phases of the as-sintered ceramics tested in the REHPTS. PLS: pressureless sintering; HP: hot pressing.  Label   Initial composition  Sintering type, temperature, time,  (vol%)  environment, pressure  Final density (g/cm3 )  Final relative  density (%)  Final main phases   Composition minor  HCM   ZCM   ZCT   HfC + 20 MoSi2 ZrC + 20 MoSi2 ZrC + 20 TaSi2  PLS 2200 K, 60 min, Ar, -  PLS 2200 K, 60 min, Ar, -  HP 1970 K, 6 min, vacuum, 30 MPa   11.1  6.2   7.1   97   95   99   HfC, MoSi2 ZrC, MoSi2 ZrC, (Zr, Ta)C   phases (<5%)  (Hf, Mo)x Siy SiC, (Mo,Zr)x Siy SiC, (Ta, Zr)x Siy  on ZrB2 /SiC composite materials during oxidation in air around 1500 K decreased when the amount of SiC in the initial material increased from 0 up to 50 vol%. This study is focused on two carbides sintered with addition of Moor Ta-disilicide developed at ISTEC-CNR by pressureless sintering or hot pressing.14-16 Preliminary studies on the oxidation behaviour at high  temperature of ZrC/MoSi2 materials  in the same solar furnace facility17 gave encouraging results, thus this composite was studied in more details and the investigation was also extended to other two UHTCs, namely HfC/MoSi2 and ZrC/TaSi2 ,  in order  to understand  the effect of changing either the carbide matrix or the sintering additive.  2. Materials and methods  Materials were processed by  ISTEC-CNR  in Faenza,  Italy and  a  summary of processing parameters  and properties  is reported  in Table 1. The following starting compositions were used:  HfC + 20 vol% MoSi2 , labelled as HCM; ZrC + 20 vol% MoSi2 , labelled as ZCM; ZrC + 20 vol% TaSi2 , labelled as ZCT.  The composites containing MoSi2 , HCM and ZCM, were pressureless sintered at 2200 K  for 60 min  in flowing Ar,15,16 while the one containing TaSi2 , ZCT, was hot pressed at 1970 K for  6 min  under  30 MPa  loading. After  sintering,  the main constituent phases are  the starting ones and minor amounts of SiC  (5 vol%)  in ZrC-composites and mixed silicides, such as  Hf-Mo-Si,  in HfC. As  it  is shown  in Table 1, all  the ceramics were  almost  fully  dense with  relative  density  above  95%. In ZCT,  the  formation of  (Zr, Ta) C  solid  solutions  and of mixed  (Zr, Ta)  silicides  in  the final microstructure confirms the  cations  exchange  and  the  dissolution-reprecipitation in presence of TaSi2 .18 Therefore,  phenomena occurring  the final volumetric composition estimated through image analysis approximately resulted as:  80 HfC + 15 MoSi2 + 5 (Hf, Mo)xSiy for HCM, 80 ZrC + 13 MoSi2 + 5 (Zr, Mo)xSiy + 2 SiC for ZCM, 85 (Zr, Ta) C + 10 (Zr, Ta)xSiy + 5 SiC for ZCT.  For all the three composites, the formation of mixed silicides is a consequence of the cations mobility occurring in the liquid phase at the sintering temperatures. The formation of SiC in ZrCbased ceramics is attributed to the reaction between silicides and residual carbon present as an impurity in the starting ZrC powder (1.5 wt%). The samples were cut into 25 mm diameter and 2 mm thickness by Electrical Discharge Machining (EDM). The  reactor used  to perform high  temperature oxidation   is  the REHPTS (Réacteur Hautes Température et Pression Solaire,  High Temperature and Pressure Solar Reactor), implemented at the  focus of  the Odeillo 5 kW  solar  furnace and  it  shown  in Fig. 1. A flat mirror (heliostat), whose position  is servo-controlled to the apparent movement of the sun, reflects the incident solar flux  to a concentrator with  faceted mirrors. A shutter enables to control  the fraction of  the concentrated solar flux delivered to the sample placed inside the reactor and therefore its surface  Fig. 1.   Image of the REHPTS reactor implemented at the focus of the 5 kW solar furnace.  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2869  ±  ±  in situ   temperature. In this set-up, the sample is placed 25 mm above the focus of the solar furnace, so that very high temperatures may be obtained on a homogeneous 10 mm diameter area in the middle of the sample, with low surface gradient as UHTCs present good −1 ). thermal conductivities, and at very fast rate (up  to 100 K s Two mirrors enable a monochromatic (5  \\u242em) optical pyrometer (Ircon, Modline Plus) to measure the surface temperature in the centre zone  (6 mm) of  the sample  through a fluorine window. The pyrometer, together with all the parts present on the optical path - window and mirrors - was calibrated on a blackbody. The accuracy of  the  temperature measurement  is going  from 1400   15 K to 2100   22 K. The oxidations were performed  in air with an atmosphere continuously renewed. Due  to  the altitude (1500 m) of  the  laboratory, the total atmospheric pressure is around 87 kPa and the oxygen partial pressure pO2 is 17 kPa. The  temperature of  the samples was maintained at a constant plateau during 20 min and a video camera was used  to  follow  the oxidation process. The interest of solar concentrated energy is that the sample can be heated from ambient to the desired temperature in a few seconds. So we can clearly analyze  the very beginning of  the oxidation at one  temperature without any disturbance due  to a too long transient step. But for full information about the mechanisms of such a complicated oxidation process, a single stop of the process after 20 min could be insufficient. A mass spectrometer (Pfeiffer Omnistar) enables in situ gas phase analysis. As explained further in the text, CO is expected to be one of the main gaseous products during oxidation, but its molar weight is the same as N2 one (m/e = 28), so it is impossible to separate the contribution of CO from the one of preponderant N2 . We  therefore mainly followed  the signal corresponding  to m/e = 44, corresponding both to CO2 and gaseous SiO. The  samples were weighted before and after oxidation  in order to assess the mass variation (converted to a mass variation −2 min −1 ) and  rate expressed  in mg cm the surfaces and crosssections were analyzed after oxidation using XRD, SEM and EDS. In order to compare the stable products formed by oxidation of  the ceramics  in air,  thermodynamic calculations were performed using the GEMINI19 software. This software is based on Gibbs energy minimization in order to define which compounds are  the most stable at  thermodynamic equilibrium, according to the initial conditions (temperature, pressure, molar fractions of reactants). The COACH database management system provides to GEMINI the requested thermodynamic data, built from the JANAF database,20 for each possible compound. Chemical kinetics and nucleation energies are not  taken  into account by GEMINI software.  3. Results  3.1. Thermodynamic calculations  Fig. 2 presents  the moles of  solid phases  and  the molar fraction of gaseous products CO, CO2 and SiO and molybdenum oxides,  from 1000  to 2200 K  for  the  three carbides as a  function of temperature at the thermodynamic equilibrium with the following initial conditions:   10 moles of HCM/ZCM/ZCT,  150 moles of air (120 moles N2 , 30 moles O2 ), Total pressure: 105 Pa.  The molar compositions of HCM, ZCM and ZCT were fixed according to the compositions of the as-sintered materials given in Section 2 and the available thermodynamic data:  HCM: 86.5 HfC and 13.5 MoSi2 , mol%; ZCM: 85.6 ZrC, 9 MoSi2 , 2.8 ZrSi2 and 2.6 SiC, mol%; ZCT: 65 ZrC, 24.1 TaC, 5 ZrSi2 and 4.9 SiC, mol%.  ϕ  −  For simplicity, the (Zr, Ta) C solid solution of ZCT has been decomposed  into ZrC and TaC phases, which allows a  rough calculation. The number of elements N (Zr or Hf, C, Ta or Mo, Si, N, O) is 6 in any of the calculations. As there cannot be more than   = 4 phases (3 solid oxides: ZrO2 or HfO2 , MoO3 or Ta2O5 , silica or zircon + 1 gaseous phase) coexisting at any temperature, the variance (v = N + 2  ϕ) is 4. The fixed three independent parameters are total pressure, initial composition of the gaseous phase (air) and  initial composition of each solid material, hence only one  thermodynamic equilibrium  is possible for each  temperature. We have chosen to use 150 moles of air in order to provide enough oxygen (i.e. 30 moles) to oxidize completely 10 moles of material and therefore to reach the thermodynamic equilibrium for all of  the possible  reactions. These proportions may  look huge.  Indeed, oxidizing 1 mole of material under 15 moles of air - or lower proportion respecting the 1:15 ratio - would give the same conclusions, we have chosen so high proportions  in order to have more significant amounts of products and clearer interpretations, neglecting fewer minor products. Due to the oxidation kinetics, we do not expect  to reach  the  thermodynamic equilibrium during experiments  that will  last only 20 min. But the thermodynamic calculations give an indication on what the final state of  the system should be, enabling us  to estimate  the stability of one potential material in extreme conditions. Looking at  the solid species formed  in  the composites containing MoSi2 ,  in Fig. 2a and c, HCM should be oxidized  into hafnia and silica at any  temperature. ZCM should form zirconia and zircon up  to 2000 K, and above  this  temperature only zirconia and silica should be stable. Molybdenum solid oxide, MoO3 , is stable only at 1000 K. Moving to the volatile species in HCM and ZCM,  the plots  in Fig. 2b and d have analogous trends:  the main gaseous product  is CO2 whose molar fraction remains between 6% and 7%,  then O2 in  the atmosphere  in both cases, showing there was enough oxygen to complete any potential oxidation. Other gaseous products are CO, SiO, and molybdenum oxides. These last have a general formula MonO3n with n = 1-5. The  increase  in  the molar fraction of MonO3n is due  to  the  fact  that higher  temperatures  favour  the  formation of “lighter” oxides, Mo3O9 being preponderant  from 1100  to 1800 K, MoO3 from 2000 K. The molar  fraction of molybdenum oxides  is a bit  lower  for ZCM  than  for HCM, which  is    \\x0c', '2870   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Fig. 2. Amount of solid phases (left) and molar fraction of   the main gaseous products (right) at   the   thermodynamic equilibrium as a function of   the   temperature,  calculated using GEMINI software: (a, b) for HCM, (c, d) for ZCM and (e, f) for ZCT.  expected as HCM  is  richer  in MoSi2 than ZCM. The  interest of using ZCT  instead of ZCM or HCM could be  the possibility  to form  tantalum solid oxide Ta2O5 at any  temperature, as Fig. 2e  shows. Consequently, besides CO2 , CO and SiO, no other gaseous oxide would evolve  in  the  investigated  temperature range (Fig. 2f). The oxygen consumption is a bit lower, due to the fact that the formation of Ta2O5 requires fewer moles of oxidant  than  the  formation of MonO3n .  In all  the  three cases, one can see that conspicuous CO escape occurs around 2000 K. From these thermodynamic calculations, it seems that ZCT is the most promising composite, due  to  the formation of more solid products and  less gaseous species  than HCM and ZCM. Since the  formation of solid phases has slower kinetic as compared to gases, ZCT is expected to undergo less degradation than the other two ceramics.  3.2.   In situ analyses  3.2.1. Video images  Fig. 3 presents video captions of  the various materials after 15 min oxidation tests in air at 1800, 2000 and 2200 K, respectively. We can observe that at 1800 K, the oxide layer that forms  ×  ×  vs. 5.8   ×  ×  −1  −1 vs. 8   on HfC is rough and broken into several fragments, whereas the oxide layer that forms on ZrC materials looks quite smooth and well adherent  to  the carbide. This could be explained by a difference  in  the  thermal expansion coefficient: HfC has a higher linear coefficient of thermal expansion than HfO2 (up to 1670 K: −6 K −6 K −1 ,  7.3   10  10 respectively),21,22 so  the oxide  layer  tends  to break due  to  the expansion of  the carbide. On  the other hand, ZrC has a  lower coefficient  than ZrO2 (up −6 K −6 K −1 , respectively),21,22 to 1670 K: 7.6   10  10 that means that the oxide could dilate more than the carbide and results in a complete coverage of the carbide. At 2000 K, HCM still appears rough and porous and has a bigger volume than the starting disc. We can also notice  that HCM has slightly moved from the support due to the release of gaseous products also on the back face. ZCM clearly shows the formation of bubbles on the top, probably due to the formation of liquid silica, and ZCT is the least damaged and has an aspect not notably different from the photo taken at 1800 K. At 2200 K, the composites containing MoSi2 , HCM and ZCM, display clear boiling phenomena, while for ZCT just a little bump can be observed on the upper surface. Therefore, according  to  these visual  inspections,  it seems  that ZrC has better oxidation resistance than HfC, and that TaSi2 has  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2871  Fig. 3. Video captures of the three UHTC after 15 min. Oxidation in air at (a) 1800 K, (b) 2000 K and (c) 2200 K.  more beneficial effect  than MoSi2 ,  in agreement with  the  thermodynamic calculations that showed that the oxidation of ZCT produces more solid oxides and less gaseous products than the ZCM and HCM.  3.2.2. Mass spectrometry  Fig. 4 displays  the evolution of CO2 and/or SiO concentration as a  function of  the oxidation  temperature  for  the  three carbides, determined using mass spectrometry. We can observe that, at 1800 and 2000 K, HCM is the material that releases more gaseous products during oxidation, in agreement with the video images (Fig. 3). This can be due to the fact that its oxide scale is subjected to continuous rupture and allows the gases escape from the oxide/carbide interface. Nevertheless, after about 10 min at 2000 K and from  this  temperature (Fig. 4b and c),  the slope of the HCM curve is less steep than at the beginning of the oxidation. The apparition of a liquid glass and its vigorous migration to  the surface  in droplet shape, as confirmed  in Fig. 3, could partially fill  the porous structure and prevent gaseous products escape. Equivalent amount of CO2 and SiO gases is detected in ZCM and ZCT during  the oxidation at 1800 K and during  the first 10 min of oxidation at 2000 K. During  the second half of the oxidation at 2000 K, ZCT evolves higher amount of gaseous products. This difference can be due to the fact that during sintering, (Zr, Ta) Si2 mixed silicides are formed and they are less stable than MoSi2 at high temperature. ZrSi2 has a melting point around 1900 K23 and has  the  tendency  to decompose around  these  temperatures, releasing more SiO. A second explanation comes  from  the video observations:  the  liquid oxide migrates to the surface of ZCM and can therefore prevent gaseous products  from being  released. Reduced amount of silicides  in  the as-sintered ZCT seems  to  limit  the  formation of big bubbles inside  the  liquid oxide,  therefore gaseous products may faster escape  through fissures and  the bursting of  smaller bubbles. Both hypotheses could also explain why at 2200 K  the gases escaping  trend changes: ZCT  releases more CO2 and/or SiO than  the  two other materials, which have now a comparable behaviour. We observe  that  the higher  is  the  temperature,  the higher is the amount of gaseous products released by ZrC-based materials.  3.3. Post-experimental characterizations  3.3.1. Mass variation rates  Table 2 gives  the global mass variation rate for each mate−2 min −1 . In general,  rial expressed  in mg cm the higher  is  the temperature,  the faster  is  the oxide growth and  the mass gain. The mass variation rate of HCM could not be evaluated owing to  important breaking and  removal of  the oxide  layer on HfC during  the experiments. The nature of  the additive  in  the ZrC materials has  little  influence at 1800 K, but at higher  temperature,  the mass variation  rate  is more  important when TaSi2 is used. Tantalum seems therefore to play a significant role in the  Table 2  Average mass variation rates and crystalline phases obtained by XRD for each material as a function of the oxidation temperature.  Mass variation rate −2 min −1 ) (mg cm  Crystalline phases by  XRD  Temperature (K)   1800   2000   2200   1800   2000   2200   HCMa  n.a.   n.a.   n.a.   ZCM   1.28   1.72   2.23   ZCT  1.20  2.37  3.86  HfO2 , HfSiO4 traces  HfO2 , HfSiO4 HfO2 , HfSiO4  ZrO2 , MoSi2 , Mo5 Si3 traces  ZrO2 ZrO2  TaZr2.75O8 , ZrO2 TaZr2.75O8 , ZrO2 TaZr2.75O8 , ZrO2  a Loss of external surface during test.  \\x0c', '2872   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  0  0.3  0.6  0.9  1.2  1.5  0  20 0  40 0  60 0  80 0  100 0  12 00  Time (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  0  0.3  0.6  0.9  1.2  1.5  0  200  400  600  800  1000  1200  T ime (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  0  0.3  0.6  0.9  1.2  1.5  0  20  0  40  0  60  0  80  0  100  0  120  0  T ime   (s)  C  O  2  a  d n  S  i  O  p  r  u d o  e c  d  (  %  m  o  . l  )  a)  b)  c)  HCM  ZCM  ZCT  Fig. 4. Measured concentration of m/e = 44   (SiO and CO2 produced) at  during  the  oxidation  of  the UHTCs  (a)  1800 K,   (b)   2000 K   and   (c)   2200 K   specimens.  0   500   1000   1500   2000   2500   3000   3500   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  2θ  (°)   2θ  (°)   2θ  (°)   t-HfSiO4   m-HfO  2   a)   0   500   1000   1500   2000   2500   3000   3500   4000   4500   5000   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  m-ZrO2   b)   0   500   1000   1500   2000   2500   3000   3500   20   30   40   50   60   70   80   I  n  t  e  n  s  i  t  y  (  c  n u o  t  s  )  m-ZrO2   o-TaZr2.75O8   c)   Fig. 5. XRD patterns of   the samples oxidized during 20 min   in air at 2000 K:  (a) HCM, (b) ZCM and (c) ZCT.  oxide  formation, especially at 2200 K, as SEM  will demonstrate.  investigations  3.3.2. X-ray diffraction  X-ray diffraction results performed on the surface of the three composites after oxidation  tests are  summarized  in Table 2. Fig. 5 compares the diffraction patterns for the samples oxidized at 2000 K, for each kind of material. The reference patterns used for the indexation of the peaks are: 83-0944 (monoclinic ZrO2 ), 65-1142 (monoclinic HfO2 ), 77-1759 (tetragonal HfSiO4 ) and 42-0060 (orthorhombic TaZr2.75O8 ). It has to be underlined that in HCM the external layer was detached and some pieces were lost during  the  tests at all  temperatures,  so  the X-ray analyses refer  to  the remaining exposed surface (Fig. 5a). The main crystalline phases  left at 1800 K  in HCM are monoclinic hafnia and traces of tetragonal hafnium silicate (hafnon), HfSiO4 . As  the oxidation  temperature  increases,  the amount of hafnon increases as well. Since hafnon decomposition  is  reported  to                                                                                                                            \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2873  occur above 2000 K24 that  is around  the  temperatures  tested in  the present work, we can explain  increasing amount of  this phase at higher temperatures probably in relation to the removal of a thicker and thicker outermost layer. In the ZCM composite (Fig. 5b), no traces of zircon (ZrSiO4 ) were detected at any temperature and  the main crystalline phases  is monoclinic ZrO2 with an evident preferential orientation along  the  (0 0 2)  lattice plane at  the highest oxidation  temperature. MoSi2 peaks, cubic ZrO2 and traces of Mo5Si3 are visible only at 1800 K. At all  temperatures,  the main crystalline phase of ZCT  is a mixed oxide with composition TaZr2.75O8 with an orthorhombic structure with preferred orientation along  the  (0 2 0)  lattice plane (Fig. 5c); its peaks become sharper while increasing the temperature,  indicating  improved crystallization. Since  the scattering coefficient for the orthorhombic TaZr2.75O8 phase has not been published  in  the  literature or  in available  ICSD database, we considered  this phase  as  a  solid  solution between 0.5 moles orthorhombic Ta2O5 (#54-514) and 3 moles  tetragonal ZrO2 (#42-1164) and estimated a scattering coefficient of 4.7. Monoclinic ZrO2 is also present and  its peaks signal  increases with temperature up to about 25 vol% at 2200 K, indicating the higher stability of pure oxide over  the mixed one. Cubic and  tetragonal ZrO2 could be present  too, but  the superimposition of  the main peaks with the mixed oxide hinders a conclusive analysis. Other authors25,26 and similar  tests on ZrC-based composites showed  indeed  that carbon, coming  from  the oxidation of  the carbide, stabilizes c-ZrO2 at low temperatures.  3.3.3. SEM and EDS analyses  Fig. 6 presents SEM  images of  the sample surfaces  in  the central  region where  temperature was measured  for  the  three composites at 1800, 2000 and 2200 K. As a rule of thumb, dark regions correspond to silica-based glass, bright phases to heavier oxides. The  residual external  surface of HCM  looks  similar after oxidation at 1800 and 2000 K, composed by porous and cracked HfO2 , with bright contrast, and HfSiO4 , with light grey contrast and smoother appearance. At 2200 K, most of the surface is based on HfO2 and HfSiO4 , but large areas are covered with dark silica-based glass  (Fig. 6a). These observations are in agreement with the video images, revealing the vigorous formation of bubbles at 2200 K, and with  the X-ray diffraction. For ZCM (Fig. 6b), at 1800 K,  the surface  is made of cracked ZrO2 and some MoSi2 phase can be observed, with an irregular shape and smooth edges. Increasing  the oxidation  temperature to 2000 K, silica migration to the surface occurs. ZrO2 phase is present in two morphologies: as cracked large grains (10  \\u242em) and brain-like smaller grains,  indicating precipitation, nucleation and growth  from  liquid phase. At  this  temperature,  large areas of the sample surface were covered with dark silica-based glass, accordingly to video images. Further temperature increase induced the formation of larger flat ZrO2 grains, deriving from the melting and recrystallization of  this phase  that  left silica at the grain boundaries. The addition of TaSi2 to a ZrC matrix generated a variety of articulate and impressive morphologies at the various oxidation temperatures (Fig. 6c). At 1800 K the surface has a rough cracked aspect with TaZr2.75O8 and ZrO2 being the  \\u242em   main phases. The surface of ZCT at 2000 K is mainly composed of petal-like grains of TaZr2.75O8 which  form volcanoes and tend to microfissuration. At 2200 K melting and recrystallization of  the mixed oxide occurred and 20  large grains precipitated  leaving  residual silica, containing Zr  traces, at  the grain boundaries. These grains are  in  turn composed of polyhedral structures of ZrO2 and TaZr2.75O8 showing  the growth planes elegantly decorated by a dendritic irregularly shaped phase identified as ZrO2 and ZrSiO4 containing small traces of Ta. In the ZrC-based composites, besides  these  tiny shapes  in ZCT oxidized at 2200 K, no  further ZrSiO4 has been clearly detected either by XRD or SEM-EDS. As a matter of  fact,  the  formation of zircon from mixed oxides starts at around 1470 K, with highest formation rate between 1770 and 1820 K.27 At higher temperatures ZrSiO4 concentration decreases again  indicating the occurrence of the decomposition reaction of zircon, reported to start at 1820 K, which accelerates with  increasing  temperatures. Considering the quenching of the samples at the end of the test, we can reasonably understand that upon cooling, the reverse reaction from ZrO2 and SiO2 to form zircon is suppressed. Fig. 7 presents  the cross-sections of HCM, ZCM and ZCT oxidized at 2000 or 2200 K. Fig. 7a-c shows  the morphology of  the  remaining HCM oxidized at 2000 K. Keeping  in mind that HCM lost the upper part, presumably composed by porous HfO2 , the remaining surface scale is composed by dense HfO2 , run through by HfSiO4 and MoSi2 oxidation products, that are Mo5Si3 phase and SiO2 glass (Fig. 7b). Moving inward (Fig. 7c), we find a mixed oxy-carbide and MoSi2 with its oxidation products.  It has  to be noticed  that, despite  the external  layer was detached, these remaining ones are crack free and well adherent to the unoxidized bulk. Given the abundant presence of hafnon, we can deduce  that  the  left surface did not exceed  the decomposition  temperature of  this compound,  reported  to be above 2000 K.24 Fig. 7d-g shows the morphology of ZCM oxidized at 2000 K. An external amorphous SiO2 phase, not homogeneously distributed  (as  result of dropleting phenomena), covers ZrO2 grains which have brain-like shape and  together  they  form a compact dense layer (Fig. 7e). Moving inward, a porous and brittle ZrO2 layer separates a dense region composed by ZrC grains surrounded by granulous Zr-C-O, MoSi2 and Mo5Si3 with SiO2 thin glassy film and droplets (Fig. 7f). As one approaches  the bulk (Fig. 7g), SiC particles are found  instead of SiO2 and  the oxidation mechanisms of matrix and MoSi2 are evident: MoSi2 is surrounded by brighter Mo5Si3 , adjacent  to  the pre-existing SiC. These islands are in turn contoured by Zr oxy-carbide which encompasses  the original ZrC grains. These observations  for ZCM and HCM are in agreement with the mechanism proposed by Shimada28 : an  intermediate oxy-carbide forms close  to  the interface with  the carbide, until  it  is completely evolved  to  the oxide phase. ZCT oxidized at 2000 K  (not shown) has a  layered structure with the outermost scale composed of TaZr2.75O8 with some  inclusions of Si-C-O phases.  In  the second  layer, TaZr2.75O8 is still  the main phase, but as secondary phase we find already SiC instead of oxides. Moving forward towards the bulk, at around 400  \\u242em from the surface, a complex mixture of (Zr, Ta)Si2 , (Zr, Ta)-oxy-carbide and SiC phases are found. The complex  (Zr, Ta)-C-O phase  is  the  result of partial oxidation  \\x0c', '2874   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  Fig. 6. SEM images of the surface of (a) HCM, (b) ZCM and (c) ZCT after 20 min. Oxidation in air at 1800, 2000 and 2200 K.  of  the starting matrix grain, composed also by (Zr, Ta) C (see Table 1). For ZCT composite, it is interesting to thoroughly analyze  the morphology of  the oxidized scale after  test at 2200 K, which marks the limit temperature between the goodness of ZCT as compared  to ZCM. Fig. 7h-m  shows  the cross-section of ZCT oxidized at 2200 K, where all  the  thickness of  the sample  resulted modified. The outermost  thick scale  is composed by a compact ZrO2 layer where 20-30  \\u242em large porosities can be  found  (Fig. 7i).  In  this  region, SiO2 droplets  (dark) containing Zr  traces are surrounded by a brighter phase,  identified as a solid solution with possible formula (Zr0.8Ta0.2 )O2 . Right beneath this layer, the mixed TaZr2.75O8 oxide is present forming a  thick dense scale  including  isolated porosity (Fig. 7l). Moving further  inward, fine grained ZrO2 with Ta  traces and SiO2 discrete phases stand above the already mentioned (Zr, Ta)-oxycarbide and SiC phases (Fig. 7m, upper part). The centre of the sample is composed of a mixed (Zr, Ta)-oxy-carbide containing lower amounts of oxygen (Fig. 7m, lower part). This articulate morphology is a sign of complex oxidation mechanisms, including melting, phase separation and reprecipitation, occurring at 2200 K.  4. Discussion  In  light of  the multiple analyses performed during and after oxidation  in  the  temperature  range 1800-2200 K on Hfand ZrC-based materials, we can deduce several observations.  4.1. Background on the oxidation behaviour  Numerous reports have been devoted to the study of the oxidation of carbides of groups  IV-VI and  the most widespread outcomes is that transition metal carbides, generally referred to as MeC, oxidize according to the following reactions:  O2 →  MeCxOy +  C1−x + MeC  MeCxOy +  C1−x +  O2 →  MeO2 +   O2−y   COy  +  (1)  (2)  These general reactions can be extended to the case of ZCT, where a solid solution  is oxidized  to  the corresponding mixed oxy-carbide and then to the oxide (3):  O2 →  Ta)CxOy →   TaZr2.75O8   Ta)C   (Zr,  (Zr,  +  (3)    \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2875  Fig. 7. SEM   images of   the polished cross-section of   (a-c)   for HCM,   (d-g)   for ZCM after 20 min. Oxidation   in air at 2000 K and   (h-m)   for ZCT after 20 min.  Oxidation in air at 2200 K.  in,25,26,28-30  According  to  the  studies  reported  oxide growth of carbides  is parabolic  indicating protective oxidation behaviour up  to 2000 K. During  the oxidation, gaseous CO  is produced as a by-product. CO generation  introduces porosity which allows diffusion via pores. Gas formed below  the oxide layer can also lift and disrupt the oxide layer, like in the case of HCM. Nevertheless,  the generated oxide  layer  is partially protective, as testified by the parabolic rate of gases evolution which is believed to control the oxide growth. A different behaviour is observed upon oxidation at higher  temperatures, as  illustrated by previous studies.30,31 At higher  temperatures,  the evolution of gases becomes less important as the oxide layer starts to sinter significantly and  the effective volume available for gaseous diffusion  is  reduced. This  implies  that  the sintering  temperatures of the surface oxides determine a change in the oxidation behaviour. The addition of significant amounts of silicides can partially alter  the oxidation behaviour, so  the oxidation  reactions involving the silicides and the other secondary phases, like SiC, should be considered as well. Reactions  (4),  (6) and  (8) are referred to as “passive oxidation reactions”, with formation of protective SiO2 glass, while  reactions  (5),  (7) and  (9) are  generally defined as “active oxidation reactions”, owing  evolution of gaseous species (SiO and CO): 5MoSi2 +  7O2 →  Mo5Si3 +  7SiO2 MoSi2 +  (5/2)O2 →  MoO3 +  2SiO  2TaSi2 +  (13/2)O2 →  Ta2O5 + 2TaSi2 +  (9/2)O2 →  Ta2O5 +  2O2 →  SiO2 +  O2 →   4SiO2   4SiO   +  +  SiC   SiC    CO2  +   SiO    CO   to   the  (4)  (5)  (6)  (7)  (8)  (9)  These general reactions can be extended to the case of mixed silicides. Passive oxidation  is expected  to occur at  temperature up to about 1800 K, while active reactions occur at higher temperatures and in low oxygen partial pressure. The surface of the samples in the present tests falls in the region of active regimes, that  is  from 1800 K on. Solid  reaction products are observed in the samples cross-sections, owing to different oxygen partial pressure and temperature gradient. In principle, oxidation of silicides should produce silica that diffuses through the surface and  \\x0c', '2876   L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878  forms a stable and continuous silica  layer. However, previous studies on  the oxidation of  these composites have shown  that even at temperatures lower than 1500 K, no protective oxidation layer was observed on the surface.15,16 This can be due to several reasons:  there  is not enough silica  to fill all  the volume expansion associated with the formation of the porous oxidation layer and/or large CO escape resulting from oxidation of the carbide can favour the dissociation of silica to gaseous SiO. Although no continuous silica layer was found on the surface, the presence of partially filled porosity  in  the cross-section can hinder  the fast diffusion of gaseous species towards the unreacted bulk.  4.2. Effect of the carbide matrix  First, we could say  that  thermodynamic calculations are  in quite good agreement with post-test analyses, confirming  the formation of the main solid products foreseen for the tabulated compounds. Here we will not consider  the formation and stability of HfSiO4 , owing  to  the uncertainty of  the real external layer. In addition,  it has  to be underlined  that  thermodynamic data of the mixed oxides formed in ZCT, TaZr2.75O8 , are not well known and  therefore could not be  included  in  the calculations. Moreover, the calculations considered only the thermodynamic equilibrium and did not  take  into consideration  the kinetic of the reactions or the energy required for creating new interfaces, which  is  the  reason why  the  solid compounds observed are slightly different  from  those predicted by  the  thermodynamic calculations. Nevertheless,  the outermost oxide  layers  involve Hf and Si elements (HCM), Zr and Si elements (ZCM), Zr, Si and Ta elements (ZCT), which is coherent with the thermodynamic results. As expected  from  these calculations, no molybdenum solid oxides appeared, but Mo5Si3 phase was  found close  to the interface with the carbide inside HCM and ZCM materials, probably resulting from incomplete oxidation of the additive. Video images evidenced quite surprising phenomena, as HfCbased ceramic resulted notably worse  than ZrC-based ones, as reported by other authors.25 Mass spectroscopy analyses also confirmed  the consistent oxidation of HCM over ZrC-based composites up  to 2000 K  involving great amount of gaseous species. However, at 2200 K,  this picture changed and HCM showed a behaviour  similar  to ZCM, probably owing  to  the achievement of equilibrium between gas production and glass boiling on the top. Of course these are just preliminary results of very complex mechanisms  that need  to be further  investigated and thoroughly confirmed.  4.3. Effect of the sintering additive  The addition of TaSi2 rather  than MoSi2 has  further consequences on  the oxidation behaviour of ZrC.  In addition,  it is  important  to underline  that  the  total amount of silicides  in the final microstructure was notably reduced from  the starting 20 vol%. Although TaSi2 has a high melting point  (2470 K), during sintering and oxidation mixed silicides (Zr, Ta)Si2 and Si-based phases with  significantly  lower melting point were formed. On  the other hand, upon oxidation,  a mixed oxide with stoichiometry TaZr2.75O8 was formed. Some authors have  studied  the effect of Ta-addition on  the oxidation behaviour of ZrB2 , where the introduction of tantalum was for increasing the glass viscosity and thus limiting the oxygen diffusivity.32 However, some authors also recognized that the melting point of this phase could be significantly  lower  than  that of zirconia33 with consequences on the high temperature stability of the oxide. This hypothesis seems to find a confirmation in the test performed on ZCT at 2200 K, where evident melting, dissociation, evaporation and re-precipitation phenomena occurred  leaving mainly ZrO2 in  the external  layer of  the final microstructure. According  to Bhattacharya et al.,34 who computed a recent phase diagram for the ZrO2-Ta2O5 system, Zr-Ta-O solution melts from 2100 K, that could explain why no deformation is observed up to 2000 K. The evident consequence is the better aspect of the ZCT samples as  compared  to  the ZCM ones  (Fig. 3),  suggesting  a more beneficial effect of  the addition of TaSi2 , compared  to MoSi2 . It  is probable  that  the petal-like open structure of ZCT and  the complex architecture of the modified layer allow for controlled gaseous  products  release,  avoiding  excessive  pressurization inside  the sample,  thus preserving  the scale from spalling and delamination. Only at 2200 K, ZCT showed some deformation owing  to  the  formation of  a dense  thick ZrO2 layer, which favoured the formation of large gas pockets below this scale.  4.4. Oxidation mechanisms in the different temperature ranges  According to the analyses performed and the studies carried out, we can now delineate the most important phenomena occurring during oxidation of these carbides in the temperature range between 1800 and 2200 K.  1800 K: At  this  relatively  low  temperature,  the  oxidation behaviour of ZrC-based composites is better than HfC, due to that fact that the ZrO2 -based oxide formed has the tendency to shrink and the effective volume available  for gaseous diffusion  is  reduced. On  the other hand, at this temperature, the oxidation of HfC is still dominated by  the escape of gases  through  the highly porous oxide layer. Gas evolution data suggest a linear rather than parabolic behaviour for all the composites. 2000 K: The oxidation of HfC/MoSi2 is still dominated by the escape of gases through the highly porous oxide layer. For the ZrC composite containing MoSi2 , data in Fig. 4 suggest a parabolic behaviour. In this case, it is likely that sintering of the ZrO2 oxide becomes important and protects the composite from diffusion of oxygen. The composite containing TaSi2 has a linear behaviour up to 600 s. After that time, there is an abrupt increase of gas production, that is probably related to the evolution of the cited low melting phases, especially (Zr, Ta)xSiy . At this temperature, sintering occurred for zirconia scale but not yet for the hafnia oxide layer. 2200 K: At this temperature, the HfC/MoSi2 composite shows a marked improved behaviour in comparison with the lower  temperatures, which could be attributed  to  the sintering of  the hafnia  scale. Data on gas evolution  \\x0c', 'L. Charpentier et al. / Journal of the European Ceramic Society 33 (2013) 2867-2878   2877  indicate  that  the extent of oxidation activity  is much reduced at 2200 K. Moreover, the parabolic behaviour recorded for  the gas evolution  is comparable or even better  than  that of ZrC/MoSi2 composite. In contrast, the ZrC/TaSi2 composite changes  its behaviour from protective  to non-protective. This  is due  to  the  instability of  the external TaZr2.75O8 phase at 2200 K and indicates that this oxide is not so effective in hindering the oxygen diffusivity. This results in enhanced incorporation and diffusion of oxygen down  to  the core of the sample, as well as release of gaseous species, with a linear increase with time.  5. Conclusions  Oxidation studies on Hfand Zr-carbide composites sintered with MoSi2 or TaSi2 were performed  in  the high  temperature range of 1800-2200 K  in order  to discriminate  the oxidation performances of the composites depending on matrix, sintering additive and temperature regime. Oxidation of HfC/20 vol% MoSi2 occurs with  formation of  gaseous  products  and  a  fragile  and  brittle HfO2 layer, accompanied by a notable volume expansion which causes  its delamination from  the bulk already at 1800 K. At 2200 K conspicuous melting of silica, deriving  from  the oxidation of  the silicide, and  its migration  to  the surface  take place with  formation of glass drops on the sintered HfO2 external scale. The same silica drops were observed  in ZrC/20 vol% MoSi2 composite from 2000 K on. At higher temperatures zirconia melting occurred with consequent limitation of gaseous products escape. Changing  the additive of  the ZrC matrix and using 20 vol% TaSi2 appears  to reduce  the deformation up  to 2000 K,  thanks to  the formation of a mixed oxide (TaZr2.75O8 ) partially filled with silica. Its peculiar open solid structure immersed in viscous glass allows the system to accommodate the new oxide structure without notable  sample distortion and, at  the  same  time,  the continuous glass hinders gases escape or oxygen penetration. Problems occur when the temperature achieves 2200 K: melting of the mixed oxide occurs, vigorous evaporation of CO and SiO is allowed and oxygen is let flowing all across the sample. The oxidation mechanisms of  these carbide composites are a very complex matter  indeed and  this study does not claim  to be conclusive, however it pointed out that there is not a material suitable for all  the  temperature conditions: ZrC + TaSi2 seems more adequate  for  temperatures up  to 2000 K, HfC + MoSi2 has better performances at  temperatures  from 2200 K on and ZrC + MoSi2 has  intermediate behaviour between  the  two. A following step could be the design of a material containing low amount of manifolds phases: TaSi2 to induce glass immiscibility  in medium-high  temperature regime and HfC for oxidation protection in ultra-high temperature regime. ZrC-based materials could be used to design future solar receivers if the working temperatures will not overpass 2000 K. HfC does not seem adequate due  to  the  formation of brittle solid particles  that could be drained  in  the  coolant  and may  erode  the  tubes  and  the turbine.  References  1. 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High temperature chemistry and oxi the  ZrC/ZrO2 interface  1996;11:2594-7.  27. Kaiser A, Lobert M, Telle R. Thermal stability of zircon. J Eur Ceram Soc 2008;28:2199-211.  28. Shimada S. Interfacial reaction on oxidation of carbides with formation of carbon. Solid State Ionics 2001;141-142:99-104.  29. Rudneva VV, Galevskii GV. Investigation of   thermal oxidation resistance  of nanopowders of  2007;48:143-7.  refractory carbides and borides. Russ J Non-Ferr Met  dation of ZrB2 ceramics containing SiC, Si3N4 , Ta5 Si3 , and TaSi2 . J Am Ceram Soc 2008;91:2250-7.  33. Levin EM, Robbins CR, McMurdie HF. Phase diagram   for   ceramists.  Columbus: The American Ceramic Soc., Inc.; 1964.  34. Bhattacharya AK,   Shklover V,   Steurer W, Witz G,   Bossmann HP,  Fabrichnaya O.   Ta2O5 -Y2O3 -ZrO2 thermodynamic  description.   system:   experimental   J   Eur Ceram   Soc   study  and 2011;31:  preliminary   249-57.  \\x0c']"
},{
  "_id": 85,
  "PDF": "High-Entropy Metal Diborides A New Class of High-Entropy Materials and a New Type of Ultrahigh Temperature Ceramics.pdf",
  "Text": "['OPEN  received: 13 July 2016  Accepted: 31 October 2016  Published: 29 November 2016  High-Entropy Metal Diborides:  A New Class of High-Entropy  Materials and a New Type of  Ultrahigh Temperature Ceramics  Joshua Gild1, Yuanyao Zhang1, Tyler Harrington1, Sicong Jiang1, Tao Hu2, Matthew C. Quinn2,  William M. Mellor2, Naixie Zhou1, Kenneth Vecchio1,2 & Jian Luo1,2  Seven equimolar, five-component, metal diborides were fabricated via high-energy ball milling and  spark plasma sintering. Six of them, including (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2) B2, (Hf0.2Zr0.2Mo0.2Nb0.2Ti0.2)B2, (Hf0.2Mo0.2Ta0.2Nb0.2Ti0.2)B2, (Mo0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, and  (Hf0.2Zr0.2Ta0.2Cr0.2Ti0.2)B2, possess virtually one solid-solution boride phase of the hexagonal AlB2  structure. Revised Hume-Rothery size-difference factors are used to rationalize the formation of highentropy solid solutions in these metal diborides. Greater than 92% of the theoretical densities have  been generally achieved with largely uniform compositions from nanoscale to microscale. Aberrationcorrected scanning transmission electron microscopy (AC STEM), with high-angle annular dark-field  and annular bright-field (HAADF and ABF) imaging and nanoscale compositional mapping, has been  conducted to confirm the formation of 2-D high-entropy metal layers, separated by rigid 2-D boron  nets, without any detectable layered segregation along the c-axis. These materials represent a new  type of ultra-high temperature ceramics (UHTCs) as well as a new class of high-entropy materials, which  not only exemplify the first high-entropy non-oxide ceramics (borides) fabricated but also possess a  unique non-cubic (hexagonal) and layered (quasi-2D) high-entropy crystal structure that markedly  differs from all those reported in prior studies. Initial property assessments show that both the hardness  and the oxidation resistance of these high-entropy metal diborides are generally higher/better than the  average performances of five individual metal diborides made by identical fabrication processing.  Recently, the fabrication and properties of metallic high entropy alloys (HEAs) have attracted significant research  interests1,2. In an HEA, the configurational entropy of a solid-solution phase is maximized to stabilize it against  the formation of intermetallics. Typically, five or more elements can be mixed in a HEA in equimolar concentrations to produce a maximum molar configurational entropy of Δ Smix =  RlnN, where N is the number of equimolar components and R is the gas constant1,2. HEAs have shown superior mechanical and physical properties1-3;  specially, a series of recent studies fabricated a class of refractory, metallic HEAs and demonstrated their excellent wear resistance and strength, including (especially) exceptional high-temperature properties4-8. Since the  minimization of Gibbs free energy (G =  H− TS, where H is enthalpy, S is entropy, and T is temperature) dictates  the thermodynamic stability of a material at a constant pressure, a high-entropy material (with large S) can be  thermodynamically more stable (particularly) at high temperatures, motivating this study to explore the phase  stability and fabrication feasibility of high-entropy metal diborides, as a new type of high-entropy materials as  well as a new class of ultra-high temperature ceramics (UHTCs). Most prior studies of crystalline high-entropy materials have been conducted for metallic HEAs of mostly  simple faceand body-centered cubic (FCC and BCC), as well as occasionally hexagonal close packing (HCP),  crystal structures1,2; much less studies have been done for making crystalline high-entropy ceramics (albeit  that glasses can be considered high-entropy materials in a broad definition), particularly those with more complex, non-cubic, crystal structures. Most recently, Rost et al. successfully fabricated an entropy-stabilized oxide,  (Mg0.2Co0.2Ni0.2Cu0.2Zn0.2)O, that possessed a single-phase rocksalt (which is also a FCC) structure when it was   1Program of Materials Science and Engineering, University of California, San Diego, La Jolla, CA 92093-0448,  USA . 2Department of NanoEng ineer ing , Un ivers ity of Ca l iforn ia , San D iego , La Jo l la , CA 92093-0448 , USA .  Correspondence and requests for materials should be addressed to J.L. (email: jluo@alum.mit.edu)  1  Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946www.nature.com/scientificreports\\x0c', 'HEB #1  HEB #2  HEB #3  HEB #4  HEB #5  HEB #6  HEB #7  Composition  (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2  (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2  (Hf0.2Zr0.2Mo0.2Nb0.2Ti0.2)B2  (Hf0.2Mo0.2Ta0.2Nb0.2Ti0.2)B2  (Mo0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2  (Hf0.2Zr0.2W0.2Mo0.2Ti0.2)B2  (Hf0.2Zr0.2Ta0.2Cr0.2Ti0.2)B2  Single Boride  Phase?  Yes  Yes  Yes  Yes  Yes  No  Yes  δa  1.4%  1.7%  1.7%  1.3%  1.6%  2.0%  2.3%  δc  3.9%  5.2%  5.2%  4.0%  4.6%  6.2%  5.2%  a (Å)  Average  3.110  3.093  3.101  3.084  3.090  3.082  3.081  c (Å)  Average  3.346  3.307  3.311  3.253  3.265  3.268  3.307  XRD  3.361  3.316  3.345  3.279  3.253  —  3.336  XRD  3.101  3.080  3.092  3.082  3.075  —  3.079  Relative  Density  92.4%  92.4%  92.3%  92.2%  92.1%  —  92.2%  Table 1.    Summary of the seven metal diboride systems studied. For the lattice parameters (a and c), the   “average” values represent the means of five individual metal diborides while the “XRD” values represent the  actual lattice parameters of the high-entropy solutions measured by XRD. See Supplementary Table S-I for  additional data.  Figure 1. Schematic illustration of the atomic structure of the high-entropy metal diborides. Here, M1, M2,   M3, M4, and M5 represent five different transition metals (selected from Zr, Hf, Ti, Ta, Nb, W, and Mo). This new  class of high-entropy materials and new type of UHTCs have a unique layered hexagonal crystal structure with  alternating rigid 2D boron nets and high-entropy 2D layers of metal cations (as essentially a class of quasi-2D  high-entropy materials), with mixed ionic and covalent (M-B) bonds between the metals and boron.  quenched from a sufficiently high temperature9; subsequent studies revealed that this entropy-stabilized oxide  and its derivatives, (Mg0.2Co0.2Ni0.2Cu0.2Zn0.2)1-x-yGayAxO (where A =  Li, Na, or K), have colossal dielectric constants10 and superionic conductivities11. To the best of our knowledge, this class of entropy-stabilized oxides and  its derivatives represent the first and only crystalline high-entropy ceramics that have been reported to date. This study fur ther extended the state of the ar t for the cr ysta l line high-entropy ceramics v ia successfu l ly synthe s iz ing a new c lass of h igh -ent ropy me t a l d ibor ide s , inc lud ing (Hf 0 .2Z r 0 .2Ta 0 .2Nb 0 .2Ti 0 .2)B 2 ,   (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2, (Hf0.2Zr0.2Mo0.2Nb0.2Ti0.2)B2, (Hf0.2Mo0.2Ta0.2Nb0.2Ti0.2)B2, (Mo0.2Zr0.2Ta0.2Nb0.2Ti0.2)  B2, and (Hf0.2Zr0.2Ta0.2Cr0.2Ti0.2)B2 (Table\\xa01). This work has greatly extended the knowledge of high-entropy materials, not only since it is the first time crystalline high-entropy non-oxide ceramics (specifically borides) have  been synthesized, but also because these high-entropy metal diborides exhibit a unique layered hexagonal crystal  structure with alternating rigid two-dimensional (2D) boron nets and high-entropy 2D layers of metal cations (as  essentially a class of quasi-2D high-entropy materials), as schematically shown in Fig.\\xa01, which distinctly differs  from any other high-entropy crystalline phases reported to date.  Results  Phase Evolution and Formation of High-Entropy Ceramic Phases.   To synthesize high-entropy  metal diborides, five commercial metal diboride powders of equimolar amounts were mixed and mechanically  alloyed via high energy ball milling (HEBM) for six hours; subsequently, the HEBM powders were compacted  into disks of 20-mm diameter and densified utilizing spark plasmas sintering (SPS) at 2000 °C for 5 minutes  under a pressure of 30 MPa. The detailed synthesis procedure was described in the “Methods” section. Seven  high-entropy metal diboride compositions were tested in this study, which are sometimes referred as HEB #1-#7  (as listed in Table\\xa01 and Supplementary Table S-I. Representative X-ray diffraction (XRD) patterns shown in  Fig.\\xa02 and Supplementary\\xa0Figs\\xa0S1-S7 illustrate the phase evolution during the HEBM and SPS fabrication process. The initial mixture of powder displayed XRD peaks for five individual metal diboride phases (although  some peaks overlap for most compositions), which broadened and merged (due to the particle size reduction  and mechanical alloying effects during HEBM); eventually, a single, high-entropy, phase of the hexagonal AlB2  structure formed after SPS at 2000 °C (Fig.\\xa02; Supplementary\\xa0Figs\\xa0S1-S7). Full-range XRD patterns of the SPS  specimens are displayed in Fig.\\xa03 (and in expanded views in Supplementary\\xa0Figs\\xa0S1-S7), where six of them, i.e.,   (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2, (Hf0.2Zr0.2Mo0.2Nb0.2Ti0.2)B2, (Hf0.2Mo0.2Ta0.2Nb0.2Ti0.2)B2,   (Mo0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, and (Hf0.2Zr0.2Ta0.2Cr0.2Ti0.2)B2, exhibit largely a single hexagonal phase, albeit the   2  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'Figure 2. XRD patterns showing the phase evolution during the HEBM and SPS fabrication processes in  (a) HEB #2 as an examplar in an expanded scale and (b) six other specimens. Only the first three peaks  of the high-entropy hexagonal AlB2 phases are shown here for figure clarity ; full-range XRD patterns (of  2θ  =  20°-100°, showing eleven XRD peaks of the high-entropy hexagonal phases) are documented in the  Supplementary\\xa0Figs\\xa0S1-S7.  Figure 3. XRD patterns of all seven specimens after SPS at 2000 °C, where the peaks of the primary   hexagonal phase are indexed. Six of seven specimens (except for HEB #6) exhibit largely a single hexagonal  phase of the AlB2 structure, albeit the presence of minor secondary (Zr, Hf )O2 (native oxides), which are  represented by the low-intensity peaks that are not indexed here the figure clarity (but indicated by the solid  dots in Supplementary\\xa0Figs\\xa0S1-S7). As the only special case, a secondary boride phase was observed in HEB  #6, with XRD peaks matching those of the (Ti1.6W2.4)B4 compound, while the major XRD peaks still represent a  hexagonal metal diboride solid-solution phase.  presence of minor secondary (Zr, Hf )O2 phases; these secondary oxide phases are represented by the low-intensity peaks that are evident in Figs\\xa02 and 3, which are not indexed in Figs\\xa02 and 3 for the figure clarity, but indicated  by the solid dots in Supplementary\\xa0Figs\\xa0S1-S7. The formation of minor amounts of secondary oxide (ZrO2 or  HfO2) phases is commonly observed in sintered ZrB2 and HfB2 specimens, which are native oxides that are difficult to remove (because of the extreme stabilities of native oxides of ZrO2 and HfO2). As the only special case, a  secondary boride phase was observed in HEB #6, (Hf0.2Zr0.2W0.2Mo0.2Ti0.2)B2, with XRD peaks matching those of  the (Ti1.6W2.4)B4 compound, while the major XRD peaks still represent a hexagonal metal diboride solid-solution  phase (Fig.\\xa03 and Supplementary\\xa0Fig.\\xa0S6).  Compositional Uniformity.   Cross-sectional scanning electron microscopy (SEM) images and the corresponding energy dispersive X-ray (EDX) spectroscopy compositional maps of three selected specimens (after   3  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'Figure 4. Cross-sectional SEM image and the corresponding EDX compositional maps of three selected  specimens after SPS, showing the formation of largely homogeneous high-entropy solid-solution phases except  for the HEB #6 shown in (c). The compositions are largely uniform albeit the presence of minor (Zr, Hf )O2 based  native oxides, e.g., in (a), and some Nb clustering in four Nb-containing specimens, e.g., in (b). The formation of  a secondary boride phase was observed only in HEB # 6, as shown in (c). Additional EDX compositional maps  (in expanded views) of all seven specimens are documented in the Supplementary\\xa0Figs\\xa0S1-S7.  SPS at 2000 °C) are shown in Fig.\\xa04 (additional EDX compositional maps of all seven specimens are documented  in the Supplementary\\xa0Figs\\xa0S1-S7). The compositions of all specimens are largely uniform, albeit the presence of  uniformly-distributed minor secondary (Zr, Hf )O2 phases (to different extents in different specimens), as well  as the (Ti1.6W2.4)B4 secondary phase in HEB #6 (only). Less than 1 at.% W (tungsten) is present in Specimens  #1-#5 and #7 as contamination from the WC-based milling media used in HEBM. EDX mapping operating at  20 kV also found micrometer-scale Nb (niobium) localization in Specimens #1 and #3-#5, with occasional Zr and  Mo clustering occurring concurrently in the same regions. This is somewhat surprising considering the fact that  NbB2 generally forms continuous solid solutions with other metal diborides12. Presumably, the Nb localization  is due to kinetic effects and can be homogenized with annealing for a prolonged time or at higher temperatures.  In general, the compositional homogeneities are largely satisfactory, as shown in Fig.\\xa04 (and in expanded views  in Supplementary\\xa0Figs\\xa0S1-S7 for all seven specimens); they are significantly more homogenous than the typical (BCC) refractory HEAs made by casting, which usually form dendrite structures with severe compositional  segregations7.  Atomic-Resolution Structural Characterization.   AC STEM HAADF and ABF imaging has been conducted to confirm the formation of uniform solid solution at nanometer and atomic scales, particularly the formation 2-D high-entropy metal layers (separated by the rigid 2-D boron nets in the (0001) basal planes) without  any significant layer-to-layer variation (or layered segregation) of different metal atoms in different (0001) planes  perpendicular to the c-axis. The STEM-ABF and STEM-HAADF images in Fig.\\xa05(a)\\xa0and\\xa0(b) show a homogeneous  solid solution phase in the HEB #2, (Hf0.2Zr0.2Ti0.2Mo0.2Ta0.2)B2. STEM ABF and HAADF images with higher  magnification showed the atomic configuration of atoms in the view of [0110] zone axis. The atomic planes (0001)  and (0110) were indicated in Fig.\\xa05(c). The mean spacing between two (0001) planes is about 3.449 Å, which is  close to 3.316 Å measured by XRD. In Fig.\\xa05(c), the metallic atoms were highlighted by red circles on (0001) plane.  Light element B can be visualized via ABF imaging. The highlighted green dots in Fig.\\xa05(c) indicated the B atoms,  which are located between two basal planes (0001). The observed atomic configuration is consistent with the unit  cell model depicted in Fig.\\xa01. The same atomic configuration and homogeneity were also observed in different  locations (Figs\\xa0S8\\xa0and\\xa0S9) and a different specimen (Fig.\\xa0S10). A careful digital image analysis (Fig.\\xa0S11) revealed  that the measured standard deviations of lattice spacings between the basal (0001) planes are only ~0.6% of the  average measured c lattice parameter or the measured variations from STEM ABF and HAADF images are  ~0.02 Å, which directly confirmed the formation 2-D high-entropy metal layers without a layered segregation of  different metal specimens in different (0001) basal planes, where these 2-D metal layers are well separated by the  rigid 2-D boron nets in between (Fig.\\xa01). Thus, these high-entropy metal diborides can be considered as (layered)  quasi-2D high-entropy materials, as schematically illustrated in Fig.\\xa01.  Nanoscale Compositional Mapping.   The compositional homogeneity at nanoscale for the HEB #2,  (Hf0.2Zr0.2Ti0.2Mo0.2Ta0.2)B2, was confirmed by EDX maps for different metallic elements. Figure\\xa06 showed that Hf,  Zr, Ta, Mo and Ti were uniformly distributed at nanoscale. No segregation or aggregation was found throughout  the scanned area. Since these compositional maps were also taken with the electron beam being parallel to the   4  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'Figure 5. Atomic-resolution STEM ABF and HAADF images of HEB #2 (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2. (a) and   (b): ABF and HAADF images at a low magnification, showing the homogeneity of the solid-solution phase.   (c) and (d): ABF and HAADF images at a higher magnification, showing atomic configuration. The electron  beam is parallel to the  [01 0] zone axis of hexagonal structure. (0001) and ( 01 0 planes are indexed in (c). The  red circles highlight the columns of transition metal atoms (Hf, Zr, Ta, Mo and Ti). The green dots indicate the B  atoms. Additional STEM images from different regions and a different specimen are documented in the  Supplementary\\xa0Figs\\xa0S8-S10; a further digital analysis of HAADF and ABF images in Supplementary\\xa0Fig.\\xa0S11  shows that the standard variations in the (0001) lattice spacings are only ~0.6% or ~0.02 Å, indicating  homogenous mixing of five metal atoms (Hf, Zr, Ta, Mo and Ti) within the 2-D metal layers in (0001) planes.  1  1  )  [0110] zone axis, they also confirmed no layered segregation along the c-axis in (0001) basal planes; thus, this is  indeed a quasi-2D high-entropy material as illustrated in Fig.\\xa01. Additional EDX mapping at a different location  was also conducted and documented in Fig.\\xa0S11.  Densification and Lattice Parameters.   In general, greater than 92% of theoretical densities has been  achieved by SPS at 2000 °C (Table\\xa01; see Supplementary Table S-I for the actual measured densities, along with  the theoretical densities calculated using the lattice parameters measured by XRD). The lattice parameters were  measured from XRD and listed in Table\\xa01. Typically, the measured lattice parameters are within < 1% of those  calculated by the rule of mixtures (Table\\xa01), which, along with the narrow XRD peaks (where the peak widths are  much narrower than the mean differences among the five peaks of individual metal diborides, as shown in Fig.\\xa02  and Supplementary\\xa0Figs\\xa0S1-S7), indicates the formation of disordered solid solutions for all high-entropy metal  diborides made in this study (consistent with the direct STEM HAADF/ABF imaging and nanoscale compositional mapping as shown in Figs\\xa05 and 6).  Hardness and Oxidation Resistance.   Initial property assessments indicated that both the hardness and  the oxidation resistance of these high-entropy metal diborides are generally greater or better than the average  performances of the individual (conventional) metal diborides made by the identical HEBM and SPS fabrication  processing. We understand that both hardness and oxidation resistance should critically depend on microstructures; the presence of porosity and oxide inclusion, as a consequence of the HEBM procedure that we adopted  for promoting the homogenization of high-entropy solid solutions, adversely affected the hardness and oxidation  resistance. To conduct a fair assessment of the relative performance of high-entropy and conventional metal diborides, we measured six single-phase high-entropy diborides, along with a controlled group of HfB2, ZrB2, TaB2,  NbB2, TiB2, and CrB2 specimens made by the identical HEBM and SPS fabrication processing using the same  processing parameters (except for CrB2; see “Methods” section for explanation). Figure\\xa07 displays the measured  hardness of six high-entropy metal diborides (with the actual measured data being listed in Supplementary Table  S-III), which are generally greater than the averages of the hardness values measured from individual metal diborides fabricated via the same route. Because MoB2 is not an equilibrium bulk phase below 1500 °C, the averages  for HEB#2-HEB#5 that contains 20% MoB2 were calculated without MoB2. Yet, it is well established that MoB2 has  a lower melting temperature and theoretical hardness than all the other metal diborides in HEB#2-HEB#5 (HfB2,  ZrB2, TaB2, NbB2, and TiB2) so that the actual averages from the “rule of mixtures,” if we could make and measure  MoB2 via the same procedure, should be even lower. Furthermore, results from an initial oxidation resistance   5  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'Figure 6. STEM-HAADF image and the corresponding EDS compositional maps for HEB #2  (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2, showing the homogeneous chemical distribution at nanoscale. These  1  compositional maps were taken when the electron beam is parallel to the  [01 0] zone axis, showing no  significant layer-to-layer variations of metal composition in different basal (0001) planes. Additional EDX  compositional maps obtained from a different region are documented in the Supplementary\\xa0Fig.\\xa0S12.  Figure 7. Measured hardness of six single-phase high-entropy metal diborides, which are generally greater  than the “rule-of-mixtures” averages of the hardness values measured from individual metal diborides  that were fabricated via the same HEBM and SPS route. Since MoB2 is not an equilibrium bulk phase below   1500 °C, the averages for HEB#2-HEB#5 were calculated without MoB2. However, MoB2 has a lower melting  temperature and theoretical hardness than all other five other metal diborides in question; thus, the actual  rule-of-mixtures averages should be even lower. It is also important to note that the hardness can be affected  by porosity and oxide inclusions so that fully-dense and oxide-free metal diborides should have greater  hardness than these measured values. We choose to compare the measured hardness values of high-entropy and  conventional metal diborides fabricated by the same method to allow a fair assessment of relative values.  measurement of these high-entropy and individual metal diborides made by the identical fabrication processing  are shown in Fig.\\xa08, with additional data and images documented in Supplementary S13-S15. Taking HEB#1  (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2 as an example (which is a good case for considering because none of its oxides is volatile in this temperature range so that the weight gains shown in Fig.\\xa08 and Fig.\\xa0S13 are easier to interpret), Figs\\xa08,  S13 and S14 show that HEB#1 performs better than most of its individual components made with the same procedure (ZrB2, TaB2, NbB2, and TiB2) except for HfB2; it certainly performs better than the “average” performance  of these five individual metal diborides. Consistently, both HEB #1 and HEB #7 maintained their shapes even at  1500 °C, while the majority of the respective individual metal diborides (except for HfB2) that were fabricated  via the same HEBM and SPS route oxidized more severely. For example, the TiB2 specimen, which represents  one most widely-used metal diboride today, pulverized completely at 1500 °C (Supplementary\\xa0Fig.\\xa0S14). Finally,  the four MoB2-containing high-entropy diborides (HEB#2-HEB#5) exhibited interesting and diverse, oxidation  behaviors because MoO3 is volatile. Despite this, some of them still perform better than many conventional metal  diborides that do not have volatile native oxides (Figs\\xa0S13\\xa0and\\xa0S15).  Discussion  The formation of (metallic) HEAs are often predicted by using the atomic-size effect (δ) and the enthalpy of  mixing (Δ Hmix) as the two main criteria1,2. The enthalpy of mixing is difficult to quantify for the current case, so  attention is focused on analyzing the atomic-size effect. The original Hume-Rothery solid-solution rule suggests  that (rsolute− rsolvent)/rsolvent ≤  15% is one of the necessary conditions for forming a binary solid solution. Following  the same concept, the average atomic-size difference (δ) can be defined for a multicomponent HEA alloy1,2, as:  δ ≡  N  ∑  i  =  1  X  \\uf8ee \\uf8ef \\uf8ef \\uf8f0 \\uf8ef  i  1  −  r  /  i  \\uf8eb \\uf8ec\\uf8ec\\uf8ec\\uf8ec \\uf8ed  N  ∑  i  =  1  X r  i  2  \\uf8f6 \\uf8f7\\uf8f7\\uf8f7\\uf8f7 \\uf8f8  \\uf8f9 \\uf8fa \\uf8fa \\uf8fb \\uf8fa  i  (1)  6  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'Figure 8. A snapshot of the relative oxidation performance of various high-entropy and individual metal  diborides fabricated and tested with the same conditions. This figure displays percentage weight gain vs.   oxidation temperature curves during annealing in flowing dry air at 1000 °C, 1100 °C, and 1200 °C (for one hour  each) sequentially for six single-phase high-entropy metal diborides [HEB #1 =  (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, HEB  #2 =  (Hf0.2Zr0.2Ta0.2Mo0.2Ti0.2)B2, HEB #3 =  (Hf0.2Zr0.2Mo0.2Nb0.2Ti0.2)B2, HEB #4 =  (Hf0.2Mo0.2Ta0.2Nb0.2Ti0.2) B2, HEB #5 =  (Mo0.2Zr0.2Ta0.2Nb0.2Ti0.2)B2, and HEB #7 =  (Hf0.2Zr0.2Ta0.2Cr0.2Ti0.2)B2], along with six individual  metal diborides fabricated via the same HEBM and SPS route. See the “Methods” section for the experimental  procedure and Supplementary\\xa0Figs\\xa0S13-S15 for additional results, including weight gain per surface  area plots, weight percentage gains at higher temperatures, and images of all specimens after oxidation at  different temperatures. In this figure (and Supplementary\\xa0Fig.\\xa0S13), solid lines represent the high-entropy  metal diborides and dashed lines represent the individual (conventional) metal diborides made by the same  fabrication route.  where r i and X i are the atomic radius and molar content, respectively, of the i-th component. Prior studies  suggested, mostly based on empirical obser vations, that a necessar y (but not sufficient) criterion for forming a single-phase (disordered) HEA is that the computed δ of the solid solution should be sufficiently small:  δ ≤  δmax ≈  4%1 or 4.3%2. By simply plugging the values of metallic or covalent radii of the metals, and the computed δ values are in the range of 3.5% to ~8% (Table S-I in the Supplementary\\xa0Material); specifically, HEB #7  has the highest δ ≈  8%; yet, it still forms single-phase, high-entropy, solid solution. In reality, metal diborides  [M2+(B−)2] form a highly anisotropic layered structure (i.e., the hexagonal AlB2 structure13), where each metal  atom donates two electrons and the M-B bonds (between the metal and B layers) have mixed ionic and covalent  characteristics (see Fig.\\xa01). Within the 2D metal layers, M-M bonds are strained significantly by the more rigid  boron net (Fig.\\xa01). Thus, none of the available (metallic, covalent or ionic) radii can effectively represent the actual  bond lengths in the metal diborides in the AlB2 structure (Fig.\\xa01)13. Alternatively, we propose to calculate the average size difference for a high-entropy metal diboride using the  lattice constants of individual metal diborides (measured lattice parameters ai and ci for the i-th MB2, as summarized in ref. 14, instead of the atomic radii of metals), as:  and  δ =  a  δ =  c  N  ∑  i  =  1  X  \\uf8ee \\uf8ef \\uf8ef \\uf8f0 \\uf8ef  i  1  −  a  i  /  \\uf8eb \\uf8ec\\uf8ec\\uf8ec\\uf8ec \\uf8ed  N  ∑  i  =  1  X a  i  i  2  \\uf8f6 \\uf8f7\\uf8f7\\uf8f7\\uf8f7 \\uf8f8  \\uf8f9 \\uf8fa \\uf8fa \\uf8fb \\uf8fa  N  ∑  i  =  1  X  \\uf8ee \\uf8ef \\uf8ef \\uf8f0 \\uf8ef  i  1  −  c  /  i  \\uf8eb \\uf8ec\\uf8ec\\uf8ec\\uf8ec \\uf8ed  N  ∑  i  =  1  X c  i  2  \\uf8f6 \\uf8f7\\uf8f7\\uf8f7\\uf8f7 \\uf8f8  \\uf8f9 \\uf8fa \\uf8fa \\uf8fb \\uf8fa  i  (2)  (3)  Subsequently, the values of δa and δc have been computed for the seven specimens and listed in Table 1 and  Supplementary Table S-I. Interestingly, the computed δa values are small (in the range of 1.3% to 2.3% for all seven  specimens) because the M-M bonds are strained by the rigid boron net (that can deform metal cations and M-M  bond lengths towards an ideal “strain-free” value dictated by stronger B-B bonds13; Fig.\\xa01). Thus, the computed δc  values may better represent the average size difference because of less constraint along the c-axis. Coincidentally,  Specimens #1-#5 and #7, for which single-phase, high-entropy, solid solutions did form, all have computed δc values in the range of 3.9% to 5.2%, whereas HEB #6, for which single-phase did not form, has the largest computed  δc value of ~6.2%. It is interesting to further note that HEB #7 (with a simple high-entropy phase) has a greater δa  but smaller δc than those of HEB #6 (with two boride phases), suggesting that a smaller δc may be more important  than a smaller δa. However, we should emphasize that small differences in lattice parameters (measured by small δa and δc)  are only one necessary, but not essential, condition for forming high-entropy solutions. A very small δ value is   7  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'certainly not a guarantee for forming a single-phase, high-entropy, solid solution. For example, the precipitation of the secondary (Ti1.6W2.4)B4 phase in HEB #6 may be related to the facts that this (Ti1.6W1.4)B4 phase is  extremely stable or WB2 is not stable by itself; further investigation is needed here to clarify the most important  reason for the precipitation of (Ti1.6W2.4)B4 in HEB #6. Moreover, the average size differences are certainly not the only factors that determine the ability to form a  single high-entropy phase. For example, it is known13 that an average lattice parameter a of ~3.04 Å would produce “strain-free” metal layers that match the rigid boron net, thereby being favored; this may also be a factor for  HEB #7 to exhibit single high-entropy phase since its average a (of ~3.081 Å) has the closest match to the ideal  strain-free value (Supplementary Table S-I; despite that this factor also favors HEB #6, where the largest δc may be  a determining factor). In addition to the several size factors discussed above, the mixing enthalpy, as well as the  valence electron concentration, may also play an important role in determining whether a single high-entropy  phase forms1,2. It is worth making a few additional notes regarding the observed phase stabilities. First, perhaps the most interesting observation is the formation of a single-phase, high-entropy solution in HEB #7, (Hf0.2Zr0.2Ti0.2Cr0.2Ta0.2) B2, despite the limited solid solubilities of CrB2 in both HfB2 and ZrB2 12,15. Second, MoB2 is believed to be metastable at room temperature, but the hexagonal MoB2 phase could be retained in the SPS specimens16; in this study,  four 20%-Mo-containing high-entropy metal diborides have been made. Third, the starting powder W2B5 (since  WB2 is not commercially available) possessed a different structure and it has limited solubilities in all diborides  except for TiB2 17-19, which can be another reason that HEB #6 did not possess a single solid-solution phase (in  addition to the largest δc of ~6.2%). It is important to emphasize that both the hardness and oxidation resistance can be affected by the microstructure, e.g., the porosity and oxide inclusions, significantly. Thus, we choose to compare the high-entropy and  individual metal diborides fabricated using the same method to allow a fair assessment of relative hardness and  oxidation resistance (even if our specimens have high levels of porosity and oxidation inclusions due to HEBM  than those fully-dense and oxide-free specimens prepared by other fabrication routes). We expect that fully-dense  and oxide-free specimens should have higher hardness and better oxidation resistance. Although the high-entropy metal diborides do appear to exhibit greater hardness and better oxidization resistance than the average performances of the individual metal diborides (provided that they are made with the same  fabrication route), perhaps a more important advantage for adopting high-entropy materials is a large compositional design space to allow tuning of properties. This will be particularly important for improving oxidation  resistance, which depends on many (often kinetic) factors; thus, there is perhaps no simple answer on whether  high-entropy metal diborides are good or bad for oxidation resistance (and some other properties). A large compositional design space is useful for designing better protective oxide scales (with additives or in composites,  which are often necessary for real applications), representing a complex material engineering problem beyond the  scope of this study. Further systematic investigation of hardness, oxidation resistance, and other properties of the  high-entropy metal diborides, which often critically depend on the microstructure and therefore the processing  optimization, is important but beyond the scope of this study that focuses on the formation, structure, microstructure, and thermodynamic stability of this new class of high-entropy materials. In summar y, this study has successfully synthesized six sing le-phase, high-entropy, metal diborides via  mechanical alloying and SPS. In general, metal diboride-based UHTCs have ultrahigh melting points, as well  as excellent thermal and electrical conductivities, hardness, and wear and oxidation resistances13,15,20-23; thus,  they have potential structural applications in extreme environments. In addition, with a unique, layered hexagonal (AlB2) crystal structure, with alternating metal and boron layers, some metal diborides also exhibit exotic  functionality, e.g., MgB2 is a well-known superconductor. While extensive future research has to be conducted to  investigate their mechanical, chemical (oxidation), and physical properties, these high-entropy metal diborides  represent a new class of UHTCs, as well as a new type of high-entropy materials that can have unique compositions and structures that differ distinctly from any other existing materials, as well as great possibilities of tailoring  their properties via an extremely-large compositional engineering space.  Methods  Synthesis of High-Entropy Metal Diboride Specimens.   To synthesize high-entropy metal diborides,  powders of HfB2, ZrB2, NbB2, TiB2, W2B5 (to substitute WB2 that is not commercially available), CrB2 (99.5%  purity ; purchased from Alfa Aesar, MA, USA), TaB2, and MoB2 (99% purity ; purchased from Goodfellow, PA,  USA) were utilized as starting materials. Appropriate amounts of five powders were utilized to fabricate specimens of each composition (with the stoichiometry being calculated on the metal basis). The seven compositions  are listed in Table\\xa01 and referred to as HEB #1 to #7 in the text. The raw powders were mechanically alloyed via  high energy ball milling (HEBM) using a Spex 8000D mill (SpexCertPrep, NJ, USA) for six hours in WC media.  To prevent overheating, the HEBM was stopped every 60 minutes to allow cooling for five minutes. The powders  were then hand ground in an agate mortar to a 325 mesh; subsequently, they were compacted into disks of 20-mm  diameter and densified utilizing spark plasma sintering (SPS, Thermal Technologies, CA, USA) in vacuum (10−2  Torr) at 2000 °C for 5 minutes under a pressure of 30 MPa, with a heating ramp rate of 100 °C/min. The inside of  the graphite die was lined with a 25μ m-thick molybdenum foil to prevent reactions between the graphite and the  diboride specimen. The molybdenum foil was then lined with a layer of 125μ m-thick graphite paper to minimize  reactions between the foil and the outer die.  Characterization. The specimens were characterized by X-ray diffraction (XRD) using a Rigaku diffractometer with Cu Kα  radiation and scanning electron microscopy (SEM) in conjunction with energy dispersive  X-ray spectroscopy (EDX). The specimen densities were measured via the Archimedes method to an accuracy  of ± 0.01 g/cm2 and the relative densities were calculated via using theoretical densities that were determined by   8  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946\\x0c', 'the ideal stoichiometry and lattice parameters measured by XRD. The atomic and nanoscale characterization was  conducted using aberration-corrected scanning transmission electron microscopy (AC STEM); STEM high-angle  annular dark-field (HAADF) images, medium-angle annular dark-field (MAADF), and annular bright-field  (ABF) images were taken by using a 200 kV STEM (ARM-200F, JEOL) equipped with a probe Cs corrector (CEOS  Gmbh), which offers an unprecedented opportunity to probe structures with a sub-Ångström resolution. For  HAADF imaging, we adopted a probe convergence angle of ~22 mrad and a detector with inner semi-angle of  > 60 mrad. The ABF images were taken with a detector of 12-23 mrad, while MAADF images were taken with  a detector of 23-50 mrad. The energy dispersion X-ray (EDX) spectroscopy was employed to map the chemical  composition at nanoscale. The TEM samples were prepared by dual-beam FIB/SEM system (Scios, FEI).  Hardness and Oxidation Measurements.   Hardness and oxidation measurements were conducted using  all six single-phase high-entropy diborides (HEB #1-#5 and #7) and six individual metal diboride benchmarking  specimens (HfB2, ZrB2, TaB2, NbB2, TiB2, and CrB2) that were made by the same HEBM and SPS fabrication  method using the same processing parameters, with one exception that CrB2 was sintered at lower temperature  of 1800 °C because its substantially lower melting (and therefore sintering) temperature. MoB2 was not examined  because it is not a thermodynamically stable phase (and will decompose to MoB and Mo2B5) below 1500 °C.  Hardness measurements were performed with a Vickers’ diamond indenter at 200 kgf/mm2 with a hold time of  15 seconds. The indentations were examined for conformation with the ASTM C1327. The indentations averaged  20-25 μ m in width during the testing. Multiple measurements were performed at different locations of each  specimen; the mean and standard deviation are reported. The density and hardness are generally uniform at  different locations for HEB specimens #1-#5 and all six individual metal diboride specimens; however, HEB #7  has a denser outside shell and less dense inner core with different average hardness values (due to the effect of  low-melting CrB2 that promotes rapid densification near the surface); thus, the hardness values were measured at  both regions and reported in Supplementary Table S-III but only the overall mean and standard deviation were  used in comparison. The oxidation experiments were conducted in a tube furnace under flowing dry air. The  specimens were annealed at 800 °C, 900 °C, 1000 °C, 1100 °C, 1200 °C, 1300 °C, 1400 °C, and 1500 °C sequentially.  Each annealing step included a one-hour isothermal oxidation at the desired temperature with a heating ramp  rate of 10 °C/min; after the isothermal annealing, the specimens were cooled in the furnace with uncontrolled  cooling rates on the order of 100 °C/min. After each annealing step, the specimens were removed from the furnace and weighted. At low annealing temperatures, specimens were weighted directly. At high temperatures (typically1300 °C and above), many specimens reacted with the alumina crucibles so that the specimens were weighted  in the crucibles to obtain the net weight gains/losses. We found the measured weights are generally accurate for  the oxidization temperatures of 1000-1200 °C (from direct weighting of specimens) and for the annealing temperatures of 1400 and 1500 °C, where the weight changes were sufficiently large to allow to be weighted accurately  in crucibles. Outside these two temperature windows, the weight gains/losses were typically on the same order of  magnitude as the measurement errors; thus, those data are not reported.  References  1. Zhang, Y., Zuo, T. T., Tang, Z., Gao, M. C., Dahmen, K. A., Liaw, P. K. & Lu, Z. P. Microstructures and properties of high-entropy  alloys. Prog. Mater. 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FirstView,  10.1557/jmr.2016.1210 (2016).  Acknowledgements  We acknowledge the financial support from an Office of Naval Research MURI program (grant No. N00014-151-2863) and we thank our Program Mangers Dr. Kenny Lipkowitz and Dr. Eric Wuchina, Principle Investigator  Prof. Dona ld Brenner, and a l l other MURI col leagues for guidance, encouragement, and helpfu l scientif ic  discussion. We thank Prof. Elizabeth J. Opila for helpful discussion about the oxidation experiments. T.H., N.Z.  and J.L. also acknowledge partial support from a Vannevar Bush Faculty Fellowship (ONR N00014-16-1-2569)  for the STEM work.  Author Contributions  J.L. conceived the idea and designed the experiments. Y.Z. conducted the initial experiments of composition  #1 before J.G. started to work on this project. J.G. conducted the most of the other experiments in a close  collaboration with of T.H. (Harrington) in the lab. All authors analyzed the data and discussed the results.  T.H. (Hu) conducted the STEM characterization and N.Z. conducted important digital analysis of the STEM  HAADF/ABF images. S.J., Y.Z., and J.G. conducted the oxidation experiments. M.C.Q. and W.M.M. conducted  the hardness measurements. J.G. and J.L. wrote the initial version of this paper; J.L., K.V., Y.Z. and T.H. revised the  manuscript critically. J.L. supervised this study.  Additional Information  Supplementary information accompanies this paper at http://www.nature.com/srep  Competing financial interests: The authors declare no competing financial interests.  How to cite this article: Gild, J. et al. High-Entropy Metal Diborides: A New Class of High-Entropy Materials  and a New Type of Ultrahigh Temperature Ceramics. Sci. Rep. 6, 37946; doi: 10.1038/srep37946 (2016).  Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and  institutional affiliations.  This work is licensed under a Creative Commons Attribution 4.0 International License. The images  or other third party material in this article are included in the article’s Creative Commons license,  unless indicated otherwise in the credit line; if the material is not included under the Creative Commons license,  users will need to obtain permission from the license holder to reproduce the material. To view a copy of this  license, visit http://creativecommons.org/licenses/by/4.0/  © The Author(s) 2016  1 0  www.nature.com/scientificreports/Scientific RepoRts | 6:37946 | DOI: 10.1038/srep37946  \\x0c\"]"
},{
  "_id": 86,
  "PDF": "High-temperature (to 1600°C) oxidation of ZrB2–MoSi2 ceramics in air.pdf",
  "Text": "['Powder Metallurgy and Metal Ceramics, Vol. 51, Nos. 1-2, May, 2012 (Russian Original Vol. 51, Jan.-Feb., 2012)   HIGH-TEMPERATURE (TO 1600\\uf0b0C) OXIDATION   OF ZrB2-MoSi2 CERAMICS IN AIR   V. O. Lavrenko,1 A. D. Panasyuk,1,2    O. M. Grigorev,1 O. V. Koroteev,1 and V. A. Kotenko1   UDC 542.943:666.3/.7   Thermal gravimetry and differential  thermal analysis are used  to study nonisothermal hightemperature oxidation of 56 wt.% ZrB2-44 wt.% MoSi2 and 86 wt.% ZrB2-14 wt.% MoSi2 ceramics  in air up to 1600\\uf0b0C. Oxide layers that form on samples at different oxidation stages are studied  using AES, EDX, and SEM analyses.  It  is concluded  that  the ZrB2-MoSi2 ceramics exhibit   exceptionally high oxidation resistance.   Keywords: ceramics, zirconium diboride, molybdenum disilicide, high-temperature oxidation, layer by-layer analysis of oxide layers.   INTRODUCTION   In recent years, various world laboratories have studied the high-temperature physicochemical properties of   ultrahigh-temperature ceramics. A major part of the publications focus on the high-temperature oxidation of   materials based on silicon carbide and refractory borides of group 4 metals [1-7]. Nevertheless, there are also   important results that deal with various aspects of the high-temperature oxidation of the most promising ceramics in  the ZrB2-MoSi2 system [8-11].    Our goal is to ascertain the kinetics and mechanism of oxidation for ceramics containing 44 and 14 wt.%  MoSi2 up to 1600°C.   EXPERIMENTAL PROCEDURE   Ceramic samples of composition 56 wt.% ZrB2-44 wt.% MoSi2 (sample 1) and 86 wt.% ZrB2-14 wt.%  MoSi2 (sample 2) were hot-pressed from powder mixtures at 1775-1800°C, with isothermal holding for 15 min,   and a pressure of 30 MPa. The samples were subjected to high-temperature oxidation in air in nonisothermal   conditions:   they were heated from room   temperature   to 1600°C at a rate of 15°C/min   in   the furnace of a   Derivatograph K 1700 (Hungary). Prior to each experiment, the furnace was purged with high-purity argon at   excess pressure for 60 min.    The starting samples and oxide layers formed on them were examined using a LAS-2000 auger   spectrometer (Riber, France). The chemical composition of samples 1 and 2 before and after oxidation at different   1Frantsevich Institute for Problems of Materials Science, National Academy of Sciences of Ukraine, Kiev,   Ukraine.     2To whom correspondence should be addressed; e-mail: panasyuk@ipms.kiev.ua.   Translated from Poroshkovaya Metallurgiya, Vol. 51, No. 1-2 (483), pp. 131-137, 2012. Original article   submitted May 29, 2009.   102                               1068-1302/12/0102-0102 \\uf0d32012 Springer Science+Business Media, Inc.     \\x0c', 'a   b  Fig. 1. DTA and TG curves for high-temperature oxidation of 56 wt.% ZrB2-44 wt.% MoSi2 (a) and  86 wt.% ZrB2-14 wt.% MoSi2 (b) ceramic sample   stages of the process was determined with energy-dispersive X-ray analysis (EDX). Electronic microphotographs of   nonoxidized samples and those oxidized at different temperatures were taken using scanning electron microscopy   (SEM).   Figure 1 shows mass change per unit surface \\uf044m / S in mg/cm2 (TG curves) and heat release (DTA curves)   for the two test materials.    Sample 1 (56 wt.% ZrB2-44 wt.% MoSi2). At   the first stage of   the process, air oxygen   is weakly   chemisorbed up to 472\\uf0b0C; adsorption surface saturation of this sample is complete at 100°C, while oxidation induced weight increment begins only at 472°C (Fig. 1a).   Another DTA peak on the oxidation curve for this sample corresponds to the oxidation of zirconium boride   at the maximum rate of 690°C:   2ZrB2 + 5O2 = 2ZrO2 (monocline) + 2B2O3\\uf0ad.   (1)   At about 950°C, the other component of the composite, molybdenum disilicide, starts oxidizing to form  more thermally stable, lower Mo5Si3 molybdenum silicide:   5MoSi2 + 7O2 = Mo5Si3 + 7SiO2.   (2)   The DTA peak of this reaction corresponds to 1050°C, and the entire process takes place in the range 950-1100°C.   In the next oxidation stage of sample 1, a SiO2 amorphous film forms on its surface; the film is stabilized  by tetragonal (cubic) ZrO2 crystallites and Mo5Si3 introduced into the SiO2 amorphous film to ensure its stability up    a   b   Fig. 2. The surface of the starting ZrB2-44 wt.% MoSi2 ceramic sample (a) and the upper film layer  on the sample oxidized at 1520\\uf0b0C (b)   103              \\x0c', 'Fig. 3. Composition of the starting sample (a), lower (b) and upper (c) oxidized layers of ZrB2- 44 wt.% MoSi2 ceramics, according to EDX   to 1500°C. The exothermic peak at   this oxidation stage occurs at 1260°C (Fig. 1a). The ZrO2 crystallites   substantially grow when they transform into a monocline modification only with increase of temperature above  1500°C. Traces of crystalline SiO2 show up as \\uf061 cristobalite at t > 1530°C. This causes some loss of very important  protective properties of the scale and leads to a weight increment of the sample from 5.0 to 7.6 mg/cm2 with   temperature increasing to 1600-1650°C. However, Auger electron spectroscopy reveals no crystalline inclusions of  ZrO2 and Mo5Si3 in the upper oxide layer on the 44 wt.% MoSi2 sample: there is only SiO2.  According to quantitative Auger electron spectroscopy, the starting 56 wt.% ZrB2-44 wt.% MoSi2 ceramic  sample contained 11.5 mol. MoSi2 per 18.7 mol. ZrB2 (Zr : B : Mo : Si = 18.7 : 42.0 : 11.5 : 27.6 at.%). Figures 2a   and 3a show electron microphotographs of starting sample 1 and results of its x-ray diffraction analysis.  Figure 3b indicates that there are 0.8 mol. ZrO2 and 0.14 mol. Mo5Si3 per 29 mol. SiO2 in the lower layer  of oxide film formed on sample 1, the lower scale layer being 45 µm thick. According to Auger analysis, Si-K\\uf061 :         : O-K\\uf061 = 27.7 : 63.3 ~ 1 : 2, Zr-L\\uf062 : O-K\\uf061 = 0.9 : 1.9 ~ 1 : 2, and Mo-L\\uf062 : Si-K\\uf061= 0.70 : 0.42 ~ 5 : 3. The boron   104       \\x0c', 'Fig. 4. Buildups of amorphous SiO2 on sample 2 oxidized at 1600\\uf0b0C   content of the scale layer decreased (through intensive evaporation of B2O3) to 0 during reaction (1), and its  molybdenum content substantially reduced through evaporation of MoO3, which forms at the first stage of MoSi2   oxidation:   2MoSi2 + 7O2 = 2MoO3 \\uf0ad + 4SiO2.  (3)  According to EDX (Fig. 3c), the upper oxide layer on sample 1 mainly represents amorphous SiО2 (10-  15 µm thick). Electron microscopy images (Figs. 2b and 4) show that this layer is inhomogeneous on sample 1   oxidized at 1520°C.  Sample 2 (86 wt.% ZrB2-14 wt.% MoSi2). Air oxygen is first chemisorbed on the surface of the sample (the   process lasts to 160°C) and then adsorption-desorption equilibrium is reached (Fig. 1b). Reaction (1) proceeds in   the range 470-745°C.  Oxidation of the other ceramic component (MoSi2) begins at 740°C by reaction (3) and then proceeds by   reaction (2) to 1150°C. In the fourth stage of interaction between sample 2 and air oxygen (in the range from 1155  to 1530°C), a stable amorphous SiO2 film forms on its surface, with inclusions of ZrO2 crystallites and Mo5Si3  silicide. At the final oxidation stage of samples 2 and 1, there is a substantial weight increment (to 5.8 mg/cm2) at   1540-1600°C.    According to EDX analysis, quantitative Auger spectroscopy, and electron microcopy of a cross-section of  sample 2 oxidized to 1500°C, the external scale layer contains amorphous SiO2 with crystalline inclusions of ZrO2  and Mo5Si3 (ZrO2 content is much greater than Mo5Si3 content).  Electron microscopy (Fig. 5a) illustrates that the content of ZrO2 is much greater (as compared with that of  SiO2) in the external scale layer. Zirconium dioxide forms in the upper scale layer as elongated acicular and   lamellar crystallites.   a   b  Fig. 5. Upper oxide layer on sample 2 (a) and its surface after oxidation at 1700\\uf0b0C (b)   105         \\x0c', 'Fig. 6. Different oxide layers of sample 2 oxidized to 1600\\uf0b0C (cross-section)   In individual experiments, we oxidized sample 2 to 1650 and 1700°C. In this case, surface buildups in the  upper part of the scale cross-section about ~7-8 µm thick (Fig. 6) consisted of amorphous SiO2, while the next,  deepest scale layer (15-20 µm thick) was amorphous SiO2 with a great amount of ZrO2 and Mo5Si3 crystalline  inclusions. The content of ZrO2 is much higher than that of Mo5Si3.   The upper part of Fig. 5b clearly shows the near-surface layer of scale, containing a great number of  acicular and lamellar ZrO2 crystallites formed on sample 2 at 1700°C, while the lower part of Fig. 5b shows the  upper layer consisting of vitreous SiO2 formed at 1700°C.   The research permits conclusions that the composition, morphology, and structure of different oxide layers  formed on ZrB2-MoSi2 ceramic samples may largely differ.    CONCLUSIONS   The ZrB2-MoSi2 ceramic materials show exceptionally high oxidation resistance, which is ascertained by   TG and DTA and examination of oxide layers formed on these materials. The oxidation of sample 2 with 14 wt.%  MoSi2 up to 1600°C resulted in weight increment \\uf044m / S = 5.7 mg/cm2 and up to 1530°C in \\uf044m / S = 3.8 mg/cm2.   Hence, this material belongs to ultrahigh-temperature ceramics.   The study has been conducted under Project R286 of the Science & Technology Center of Ukraine.   ACKNOWLEDGEMENTS   REFERENCES   W.-B. Han, X.-H. Zhang, J.-C. Han, and S.-H. Meng, “High-temperature oxidation at 1900\\uf0b0C of ZrB2-  xSiC ultrahigh temperature ceramic composites,” J. Am. Ceram. Soc., 91, No. 10, 3328-3334 (2008).  O. N. Grigoriev, B. A. Galanov, V. A. Lavrenko, et al., “Oxidation of ZrB2-SiC-ZrSi2 ceramics   in   oxygen,” J. Europ. Ceram. Soc., 30, No. 11, 2397-2405 (2010).   Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Evolution of structure during oxidation of zirconium  diboride-silicon carbide in air up to 1500\\uf0b0C,” J. Europ. Ceram. Soc., 27, 2495-2501 (2007).  P. Lespade, N. Richer, and P. Yoursat, “Oxidation resistance of HfB2-SiC composites for protection of   carbon-based materials,” Acta Astronautica, 60, 858-864 (2007).   F. Monteverde and A. Bellosi, “Development and characterization of metal-diboride-base composites   toughened with ultrafine SiC particulates,” Sol. State Sci., 7, 622-630 (2005).   S. Zhu, W. G. Fahrenholtz, and G. E. Hilmas, “Influence of silicon carbide particle size on microstructure   and mechanical properties of zirconium diboride-silicon carbide ceramics,” J. Europ. Ceram. Soc., 27,   2077-2083 (2007).   1.   2.   3.   4.   5.   6.   106       \\x0c', '7.   8.   9.   10.   11.   G. Talmy, J. A. Zaykoski, and M. M. Opeka, “High-temperature chemistry and oxidation of ZrB2 ceramics  containing SiC, Si3N4, Ta5Si3, and TaSi2,” J. Am. Ceram. Soc., 91, No. 7, 2250-2257 (2008).   D. Sciti, M. Brach, and A. Bellosi, “Long-term oxidation behavior and mechanical strength degradation of  pressurelessly sintered ZrB2-MoSi2 ceramics,” Scr. Mater., 53, 1297-1302 (2005).  K. Hansson, M. Halvarsson, J. E. Tang, et al., “Oxidation behavior of a MoSi2-based composite in different   atmospheres in the low temperature range (400-550°C),” J. Europ. Ceram. Soc., 24, No. 13, 3559-3573   (2004).   K. Hansson, J. E. Tang, M. Halvarsson, et al., “The beneficial effect of water vapor on the oxidation at 600  and 700\\uf0b0C of a MoSi2-based composite,” J. Europ. Ceram. Soc., 25, No. 1, 1-11 (2005).   D. Chyrkin, V. A. Lavrenko, and V. N. Talash, “High-temperature and electrochemical oxidation of   transition metal silicides,” Powder Metall. Met. Ceram., 48, No. 5-6, 330-345 (2009).   107       \\x0c']"
},{
  "_id": 87,
  "PDF": "High-Temperature Chemistry and Oxidation of ZrB2Ceramics Containing SiC, Si3N4, Ta5Si3, and TaSi2.pdf",
  "Text": "['High-Temperature Chemistry and Oxidation of ZrB2 Ceramics  Containing SiC, Si3N4, Ta5Si3, and TaSi2  Inna G. Talmy,  w  James A. Zaykoski, and Mark M. Opeka  Naval Surface Warfare Center Carderock Division (NSWCCD), West Bethesda, Maryland 20817-5700  The effect of Si3N4, Ta5Si3, and TaSi2 additions on the oxidation behavior of ZrB2 was characterized at 12001-15001C and compared with both ZrB2 and ZrB2/SiC. Significantly improved oxidation resistance of all Si-containing compositions relative to ZrB2 was a result of borosilicate glass during exposure to the oxidizing environment.  the  formation of a protective  layer of  Oxidation resistance of  the Si3N4-modified ceramics increased with increasing Si3N4 content and was further improved by the addition of Cr and Ta diborides. Chromium and tantalum oxides  induced phase separation in the borosilicate glass, which lead to  an  increase  in  liquidus  temperature  and  viscosity  and  to  a  decrease in oxygen diffusivity and of boria evaporation from  the glass. All tantalum silicide-containing compositions demon strated phase  separation in the borosilicate glass and higher  oxidation resistance than pure ZrB2, with the effect with temperature. The most oxidation-resistant ceramics contained 15 vol% Ta5Si3, 30 vol% TaSi2, 35 vol% Si3N4, or 20 vol% Si3N4 with 10 mol% CrB2. These materials exceeded the the ZrB2/SiC ceramics below 13001- oxidation resistance of 14001C. However, the ZrB2/SiC ceramics showed slightly superior oxidation resistance at 15001C.  increasing  I.  Introduction  THE unique combination of properties (high melting temperature, high hardness and strength, good oxidation resistance, and high thermal conductivity) makes ceramics based on  ZrB2 and HfB2 promising candidates for high-temperature applications1-7 such as hypersonic vehicle aerosurfaces (leading  edges and nosecaps), which require oxidation resistance at temperatures above 20001C.  The good oxidation resistance of pure Zr and Hf diborides is  a result of the formation of an oxidation product consisting of ZrO2 (HfO2) and B2O3. Below 12001C liquid boria fills all of the porosity in the fine-grained oxides, providing oxidation protection. However, above 12001C the diborides oxidize rapidly as a  result of the evaporation of the protective B2O3 from the oxide scale. It should be noted that a fraction of boria (about 15 vol%) temperatures up to 15001C  remains within the oxide scale at  because of the high surface tension of the porous, fine-grained ZrO2 layer.8 The presence of liquid boria explains the superior oxidation resistance of these borides compared with the corre sponding carbides and nitrides, which form single-phase, porous  ZrO2 products.  scales as a result of  the formation of gaseous oxidation  In general,  the oxidation behavior of non-oxide  ceramics  largely depends on the chemical composition and properties of  the oxidation products and on the combination of physical and  chemical processes taking place on the surface exposed to oxy gen-containing atmosphere. Modification of the chemical com position of the oxide layer, leading to decreased inward diffusion  of oxygen, is an effective way of controlling oxidation resistance  of nonoxide ceramics. This modification can be accomplished by  changing the bulk composition of ceramics.  Numerous studies dedicated to the improvement of the oxidation resistance of ZrB2 and HfB2 2,5,6,9-15 resulted in the selection of ceramics with 20-30 vol% SiC as optimal, preferred  compositions. The SiC-containing ZrB2 ceramics showed relatively high oxidation resistance up to 15001C as a result of the  formation of  the protective surface layer of borosilicate glass.  Compared with liquid B2O3, melting temperature, higher viscosity, and lower oxygen diff the borosilicate glass has higher  usivity,  thus providing much more effective oxidation protec tion. Theoretical aspects of the oxidation behavior of the zirconium boride-based materials have been reported earlier.1  The increase in the oxidation resistance of ZrB2/25 vol% SiC ceramics was accomplished by the addition of CrB2, TiB2, TaB2, NbB2, and VB2 (as substitution for ZrB2),16 despite the fact that the oxidation resistance of all the modifying diborides (alone) is much worse than that of the ZrB2 and HfB2.17,18 This improvement in the oxidation resistance was related to the presence of  transition metal oxides (as a result of the oxidation of the cor responding diborides) in the borosilicate glass inducing its phase  separation (immiscibility). Both increased liquidus temperatures  and viscosities, which are characteristic features of  immiscible  glasses, and are beneficial for decreasing oxygen diffusivity and  suppressing boria evaporation from the glass.  The effectiveness of oxides in enhancing immiscibility increases with increasing metallic element cation field strength, z/r,2 where z is the valence and r is the ionic radius.19-22 Because the  cation field strengths of Ti, Nb, Ta, Mo, Cr, and V are higher  than that of Zr, these elements are effective in promoting phase  separation of the ZrO2-containing borosilicate glass. Studies on the effect of alternative sources of Si  (instead of  SiC), such as Si3N4 and silicides, on the oxidation behavior of ZrB2 ceramics are not systematic and are sometimes contradictory. Bellosi and Monteverde23 reported an increase in the 13501 and 16001C  oxidation resistance of ZrB2 with the addition of 5 vol% Si3N4. During hot pressing at 17001C, Si3N4 decomposed with the formation of BN, ZrO2 and a B-N-O-Zr-Si glassy phase. Unlike the previous data, Sato al.24 et did not observe any chemical interaction between Si3N4 fibers and the ZrB2 matrix after hot pressing at 14001C. The effect of the disilicides of Ta, Nb, W, Mo, and Zr, as well  ceramics  at  as of Zr5Si3 on the oxidation resistance of ZrB2 was studied by Shaffer.25,26 In these studies, ZrB2 ceramics containing 5-15 mol% MoSi2 showed superior oxidation performance at both 14001 and 19501C. Shaffer also suggested the addition of ZrB2/ 5-50 mol% TaSi2 mixture as oxidation-protective coatings for Mo, borides, and carbides. An increase in the oxidation resistance of ZrB2 up to 12001C with the addition of ZrSi2 was reported by Lavrenko et al.27 The authors also found solid sol ubility between the components after pressureless  sintering in  vacuum (no al.28  sintering  temperature  was  given).  Panasyuk  et  reported an improvement  in oxidation resistance of  ZrB2 with the addition of 15-50 wt% CrSi2. The experiments were conducted at 6001-12001C and 740 torr (98.6 kPa) oxygen  N. Jacobson—contributing editor  w  Author to whom correspondence should be addressed. E-mail:  inna.talmy@navy.mil  Manuscript No. 24005. Received November 19, 2007; approved February 8, 2008.  Journal  J. Am. Ceram. Soc., 91 [7] 2250 - 2257 (2008)  DOI: 10.1111/j.1551-2916.2008.02420.x  r 2008 The American Ceramic Society  No claims to original US government works  2250  \\x0c', 'July 2008  High-Temperature Chemistry and Oxidation of ZrB2-Based Ceramics  2251  pressure for 120 min. The improvement in oxidation resistance  X-ray and microscopic evaluations. XRD and SEM/EDS were  was attributed to the formation of protective ZrSiO4. Pastor and Meyer29 evaluated the oxidation resistance of  ZrB2 with additions of MSi2 or M5Si3, where M is a transition metal Zr, Ta, Cr, Mo, or W. The authors observed significant  interactions or solid solubility in the systems of ZrB2 with Zr, Cr, and Mo silicides and very little or no interactions in the  systems with Ta and W silicides. After oxidation testing for up to 100 h at 12001 and 14001C, the ZrB2/15 wt% CrSi2 composition was the most oxidation resistant. Opila et al.30 reported  that the addition of 20 vol% TaSi2 to ZrB2/20 vol% SiC improved the oxidation resistance of the material at 16271C in  air. The improved behavior was attributed to the presence of Ta,  not to the increase in Si content as a result of the introduction of  the second Si-containing component. The processing, chemical  interactions, and properties of ceramics in the whole ZrB2/Ta5 Si3 system were reported by Talmy et al.31 The purpose of the present work is to characterize the oxi dation behavior of ZrB2 modified with Si3N4, Ta5Si3, and TaSi2 as sources of Si and to compare the results with the oxidation  behavior of SiC-containing ZrB2. The introduction of Ta silicides was of particular interest because it simultaneously modifies  the oxidation layer with SiO2 to form borosilicate glass and with Ta2O5 to induce glass immiscibility for potential further improvement in oxidation resistance.  The  experimental  compositions ZrB2/5, Si3N4, ZrB2/8, 15, and 31 vol% Ta5Si3, and ZrB2/10, 20, and 30 vol% TaSi2 were selected to keep the B2O3/SiO2 ratio in the surface borosilicate glass between 1.5 and 3, assuming that ZrB2 and Si-containing additives are totally oxidized. (This ratio in  vol%  20,  and 35  the ZrB2/25 vol% SiC composition is 2.) According to the SiO2- B2O3 phase diagram,32 the melting temperature of borosilicate glass of these compositions would be about 6001-7001C. It  should be noted that SiO2 content temperatures will gradually increase as B2O3 increasing oxidation temperature and/or time.  in glass and glass melting  evaporates with  II.  Experimental Procedure  The starting borides and silicides (\\x00325 mesh and 99.5% purity) were purchased from Cerac Inc. (Milwaukee, WI). Silicon nitride (LC12, particle size 0.6 mm) was purchased from HC Starck,  Inc. (Newton, MA). The powders were mixed with a mortar and  pestle  in acetone. The mixing procedure was  repeated three  times with intermediate drying. The materials were densified by  hot pressing (Model 912G, Thermal Technology Inc., Santa  Rosa, CA) in a graphite mold in He atmosphere. The hot-press ing conditions will be given in the sections corresponding to the  specific systems.  The samples were heated to the maximum temperature with a ramp rate of 401C/min under a pressure of 5 MPa, and the full  pressure was applied when the desired densification temperature  was  reached. The hot-pressed samples were  characterized by  density and open porosity (Archimedes principle), phase com position (X-ray diffraction (XRD), Siemens Theta/Theta, Mod el D 500, Bruker AXS, Madison, WI),  and microstructure  (SEM, Model  ISI ABT SR-50A, Withington, Manchester,  UK). Energy-dispersive spectroscopy (EDS) was also used to  identify the elemental composition of phases. The boron content  data were inconclusive and are not reported in the paper.  The oxidation behavior was characterized by measuring mass changes of the bars (average size about 3 mm \\x02 4 mm \\x02 5 mm and surface area about 1.0 cm2) during furnace heating in air at 10001-15001C for 2 h. The bars were supported by their ends on  alumina semirings for maximum contact with the atmosphere.  The mass with and without the semi-rings was recorded to allow  posttest mass measurements in the event that the glassy oxida tion product reacted with the alumina. The mass change results  were an average from three to five samples. Air quenching of the  samples after furnace heating was conducted to retain the high temperature condition of the surface layer for room temperature  used to characterize the phase composition and microstructure  of the surface and cross-section of the samples after oxidation.  III.  Results and Discussion  (1)  ZrB2-Si3N4 Ceramics  The ZrB2 ceramics containing 5-35 vol% Si3N4 were prepared by hot pressing mixtures of the end-member powders at 1850 1C  and 20 MPa in He for 1 h. Unlike SiC, which is chemically compatible with ZrB2,33 Si3N4 was decomposed during hot pressing. The XRD analysis showed the presence of BN,  ZrSi2, and ZrN in addition to the residual ZrB2. The coexistence of ZrSi2 and ZrB2 observed in the present research confirms the Brukl phase equilibria data,34 but contradicts the solid solubility data reported by Lavrenko et al.27 and Kieffer and Benesovsky.35  The relative densities of  the ceramics were not determined  because of the unknown quantitative phase composition of the  final materials. However,  the samples did not have any open  porosity. The microstructure of materials containing up to 20  vol% Si3N4 in the starting mixtures consists of coarse crystals surrounded by a fine-grained phase (Fig. 1). The EDS analysis  showed only Zr in the coarse crystals (presumably ZrB2) and Zr, Si, and N in the fine-grained phase (ZrB2, BN, ZrSi2, and ZrN reaction products). The sizes of both ZrB2 and fine-phase crystals decrease with increasing Si3N4 content, and ceramics containing 35 vol% Si3N4 had a uniform fine-grained structure. The oxidation resistance of the Si3N4-containing ceramics during furnace heating increased (butmass gain decreased) with  increasing Si3N4 (Fig. 2). The lower mass gain for pure ZrB2 compared with the sample containing 5 vol% Si3N4 is due to the significant evaporation of B2O3 from pure ZrB2 above 12001C, decreasing the total weight gain from the formation of ZrO2 and B2O3 during oxidation. The significantly larger mass gain measured for the 5 vol% Si3N4-containing sample in comparison with that for pure ZrB2 resulted from the lower evaporation of B2O3 from borosilicate glass formed on the surface of this sample.  The smaller thickness of the oxidation layer for the 5 vol% Si3N4-containing sample (140 mm) compared with that for pure ZrB2 (160 mm) after 2 h oxidation at 13001C affirms improved oxidation resistance of ZrB2 with the addition of even small amounts of Si3N4. The data on the formation of glass and improvement in the oxidation resistance of ZrB2 by the addition of 5 vol% Si3N4 support the results reported by Bellosi and Monteverde.23 The thickness of the oxidized layer for the ZrB2/20 vol% Si3N4 ceramics tested at 13001C was 60 mm. Only ZrO2 was identified on the surface of all oxidized Si3N4-containing samples. By analogy with ZrB2/SiC materials, CrB2 and TaB2 were added to the ZrB2/Si3N4 ceramics to induce phase separation in  the surface  Fig. 1.  SEM of ZrB2/20 vol% Si3N4 ceramics hot-pressed at 18501C and 20 MPa for 1 h.  \\x0c', '2252  Journal of the American Ceramic Society—Talmy et al.  Vol. 91, No. 7  300  200  100  Scale Thickness 60 µm 1 140 µm 2 160 µm 3  2  3  1  5 v% Si3N4  ZrB2  20 v% Si3N4  35 v% Si3N4  ZrB2/25 v% SiC  )  2  m  /  g  (  e g n a h  C  s s  a  M  0  1200  1300  1400 1500 Temperature (°C)  Fig. 2. Mass changes during furnace oxidation (2 h) of ZrB2/Si3N4 ceramics at 12001-15001C.  )  2  m  /  g  (  s  e g n a h  C  s s  a  M  200  150  100  50  0  Baseline (ZrB2 / 20 v%Si3N4)  10 mole% CrB2  10 mole% TaB2  ZrB2 / 25 v% SiC  1200  1300 1400 Temperature, °C  1500  Fig. 3. Mass changes during furnace oxidation (2 h) of ZrB2/20 vol% Si3N4 ceramics modified with 10 mol% CrB2 and TaB2 at 12001-15001C.  the surface glass during the oxidation test. The additives were  introduced as a 10 mol% substitution for ZrB2 with constant 20 vol% Si3N4. No significant changes in the microstructure were observed for the modified ceramics compared with the baseline  material. As previously discussed, Si3N4 reacted with ZrB2 during hot pressing with the formation of BN, ZrSi2, and ZrN, which were localized in the fine-grained areas. By EDS analysis,  the modifying elements were detected only in the fine-grained  phase. The dissolution of CrB2 and TaB2 in ZrB2 was determined by XRD. The ZrB2 peaks were shifted to higher angles, indicating the contraction of the ZrB2 lattice as a result of the incorporation of CrB2 and TaB2, which have smaller lattice constants than ZrB2. The oxidation behavior of the modified compositions togeth er with that of ZrB2 with 20 vol% Si3N4 and ZrB2 with 25 vol% SiC ceramics is shown in Fig. 3. The modifying diborides  increased  the  oxidation  resistance  of  the  baseline  ceramics.  The mass gain of the ceramics containing TaB2 and CrB2 was 26 and 34 g/m2, respectively, compared with 49 g/m2 for the baseline composition after 2 h oxidation at 13001C. Higher ox idation resistance of these ceramics at temperatures below 14001C compared with that of ZrB2/SiC ceramics can be attributed to the oxidation rate difference between ZrSi2 (the Si source in these ceramics) and SiC. It is known that SiC does not oxidize  appreciably at these lower temperatures to participate in the formation of borosilicate glass.16 However, ZrSi2 oxidizes more readily and enhances the low-temperature formation of the protective borosilicate glass. Lavrenko’s observation27 of  the formation of a vitreous borosilicate film after oxidation of the ZrB2/50 wt% ZrSi2 material at 7001-8501C is consistent with the present results.  The presence of Ta and Cr oxides, which induce phase sep aration in the glass,  further improves glass-protective capabili ties. The effect of the additives was very similar at temperatures below 14001C. Above 14001C the best oxidation resistance was  shown by the CrB2-containing material. Extensive phase separation (mostly ‘‘cabbage’’-like configu rations) was observed on the surface of all the modified samples  after oxidation tests. Figure 4 shows the SEM micrograph of the surface of the CrB2-containing sample after oxidation at 13001C for 2 h. The overall XRD of the surface of this sample identified ZrO2, Cr2O3, and CrO2. The presence of Cr in the Cr41 oxidation state having a high cation field strength (1322 compared with 793 nm\\x002 for Cr31) contributed significantly to the phase  Zr - 6.6; Si - 56.8; Cr - 0 at. %  Zr - 1.9; Si - 63.9; Cr - 5.9 at. %,  Zr - 17.5;  Si - 8.9; Cr - 0  at. %  100 µm  Fig. 4.  SEM of the surface of ZrB2/20vol% Si3N4/10 mol% CrB2 ceramics after oxidation at 13001C for 2 h.          \\x0c', 'July 2008  Table I.  High-Temperature Chemistry and Oxidation of ZrB2-Based Ceramics  2253  Phase Composition of ZrB2/Ta5Si3 Ceramics After Hot Pressing at 19001C for 30 min  Ta5Si3 (tetragonal) content in  ZrB2/Ta5Si3 starting mixtures  Phases determined by XRD  Mol%  Vol%  ZrB2  ZrB2 (ss)  TaB  hexagonal) (ss)  Ta5Si3 (D88  0  2  4  10  0  8  15  31  x  x  x  x  x  x  x  x  x  XRD, X-ray diffraction.  separation in the glass and, subsequently, to the improvement in  its  oxidation protection capability, which is especially protemperatures below 14001C. The EDS analysis of  nounced at  this sample showed a significant concentration of Zr in the top  area of the ‘‘cabbage’’ (17.5 at.% compared with 1.9 at.% in the  glassy area) and the presence of Cr only in the glass.  (2)  ZrB2-Ta5Si3 Ceramics  ZrB2 ceramics modified with 8, 15, and 31 vol% (2, 4, and 10 mol%) Ta5Si3 were prepared by hot pressing starting mixtures at 1900 1C and 20 MPa for 30 min in a graphite mold and He  atmosphere. The XRD analysis revealed significant  interaction  between ZrB2 and Ta5Si3 during hot pressing. The polymorphism of Ta5Si3, as well as the chemical interactions and solid solubility in the entire ZrB2-Ta5Si3 system, is described elsewhere.31 The phase compositions of the ceramics containing 2,  4, and 10 mol% are summarized in Table I.  The ceramics containing 2 and 4 mol% Ta5Si3 consisted of stoichiometric ZrB2 and ZrB2-based solid solution with a contraction of the ZrB2 unit cell. The unit cell volume of ZrB2 decreased linearly from 30.70 to 29.89 A˚ 3 as Ta5Si3 content increased from 0 to 4 mol%. Based on the insolubility of Si and ZrSix in ZrB2 reported in the literature34,36 and the contraction observed in the ZrB2 lattice, it was deduced that Ta entered the ZrB2 cell, causing the contraction (the atomic radii of Zr and Ta are 1.60 and 1.46 A˚ , respectively).  The samples containing 10 mol% Ta5Si3 also showed the presence of ZrB2 and ZrB2-based solid solution. However, the solidsolution unit cell volume only slightly changed with an increase in  Ta5Si3 content from 4 to 10 mol%. This indicated that the solubility limit of Ta in the ZrB2 lattice, with the present reactant ratios, was achieved in the material containing about 4 mol%  Ta5Si3. The simultaneous presence of ZrB2 and ZrB2-based solid solutions is probably the result of insufficient time (30 min) dur ing hot pressing to completely homogenize the composition.  the  Additionally,  samples with 10 mol% Ta5Si3 substantial amounts of two new phases: TaB and a Ta5Si3-based hexagonal (D88) solid solution with considerable expansion of the silicide cell volume from 252.77 A˚ 3 (for pure hexagonal Ta5Si3) to 273.83 A˚ 3 for the solid solution formed in the material, as determined by XRD. The possible reaction between the  contained  components could be presented as a hypothetical equation:  Ta5Si3 þ ZrB2 ! ðZr; TaÞB2 þ TaB þ ðTa; ZrÞ5Si3B  This reaction involved the decomposition of some fraction of  the ZrB2, which provided B for the formation of TaB, as well as Zr and B for the formation of the D88-based solid solution. Both the interstitial incorporation of B and the substitution of Ta with  Zr resulted in significant D88 lattice expansion. It is also noteworthy that the TaB lattice parameters corresponded to the sto ichiometric composition, indicating no dissolution of Zr in TaB.  The observed chemical  interaction and solid solubility be tween Ta5Si3 and ZrB2 contradict the only existing literature data,29 which reports that very little or no interaction is observed even after heating at 22001C and 10 \\x002 torr (1.33 Pa) for 1 h. This dramatic discrepancy in these two studies is not easy to resolve.  the samples were 6.61, 7.18, and 8.27  8,  31  The bulk densities of g/cm3  for  vol% Ta5Si3, respectively, with close-to-zero open porosity. (The relative den containing  ceramics  and  15,  sities of  the ceramics were not determined because of  the un known quantitative phase composition of  the final ceramics.)  Evidently, a decrease in the hot pressing temperature of the Ta5Si3-containing materials compared with pure ZrB2 (22001C) can be related to chemical interactions and solid solubility be tween the components.  The microstructure of  the hot-pressed ZrB2/15 vol% Ta5Si3 sample (Fig. 5) is composed of large crystals (average size about 5-10 mm) containing predominantly Zr (by EDS) surrounded by  10 µm  Fig. 5.  SEM of ZrB2/15 vol%/Ta5Si3 ceramics HP at 19001C.  \\x0c', 'fine-grained crystals (about 1 mm) consisting of both Zr and Ta,  which is an indication that fine ZrB2 particles preferentially participated in the chemical interactions between the components.  The location of Si was difficult to identify because of the over lapping of Ta and Si peaks, but  it can be assumed that  it was  localized in the fine-grained phase.  Figure 6 shows the mass gain for ZrB2-Ta5Si3 ceramics during 2 h oxidation at 11001-14001C. Materials containing Ta5Si3 exhibited increased oxidation resistance, which is further sup ported by measurements of the thickness of after oxidation at 14001C. The thickness of the oxidized layer was 115 mm for ceramics containing 15 vol% Ta5Si3 compared with 300 mm for pure ZrB2. The presence of ZrO2 and TaZr2.75O8 was determined by XRD on the surface of oxidized samples. No crystalline siliconor boron-containing  the oxidized layer  the  compounds were identified with certainty.  A periodic pattern of glassy and crystalline  areas on the  oxidized  surface  of  the  ceramics  containing  8  vol% Ta5Si3 (Fig. 7) implies glass phase separation, which was induced by the presence of Ta51. In spite of  the overlapping of major Ta  and Si peaks, EDS analysis of the surface leads to the conclusion  that  the glassy area contains a very small amount of Zr and  practically no Ta (from the size of the ‘‘far’’, \\x19 8 keV, Ta peak), while the crystallized area contains Ta and a significant amount  of Zr. The strong peak belonging to both Si and Ta in the EDS  scan of  the glassy area can be  therefore associated predomi nantly with Si, and the peak from the crystallized area of  the  scan is from a combination of both elements.  Thus,  the improved oxidation resistance of  the Ta5Si3-containing compositions compared with pure ZrB2 is due to the formation of an immiscible borosilicate glass during oxidation.  The test results of this study did not match the results of Pastor and Meyer,29 who did not observe any improvement  in the  oxidation resistance of ZrB2 by the addition of 15 wt% Ta5Si3 (the composition corresponds to 8 vol% Ta5Si3 in the present study) after oxidation in air at 12001C for 100 h.  In spite of  the improvement  in oxidation resistance of ZrB2 with the addition of up to 31 vol% Ta5Si3, ZrB2 ceramics containing 25 vol% SiC have superior oxidation resistance at tem13001C, as  peratures  above  shown in Fig.  6. The  enhanced  oxidation resistance of Ta5Si3-containing ceramics below 13001C can be attributed to the insufficient oxidation of SiC  at these temperatures to form a continuous, protective borosil icate glass, while significant oxidation of Ta5Si3 occurred below 11001C.31  (3)  ZrB2-TaSi2 Ceramics  ZrB2 ceramics containing 10, 20, and 30 vol% (approximately 7, 15, and 23 mol%, respectively) TaSi2 were prepared by hot pressing at 20001C and 20 MPa for 20 min in a He atmosphere.  The bulk densities of the samples were 5.63, 6.44, and 6.73 g/cm3, respectively. The XRD patterns of the ceramics showed  solid solubility in the system (Fig. 8). The ceramics containing 10  vol% TaSi2 did not show the diffraction peaks of TaSi2. As the amount of TaSi2 increased to 20 and 30 vol%, the TaSi2 peaks became evident and increased in intensity with increasing TaSi2 concentration. The peaks were shifted to the left from the  position of  the  stoichiometric TaSi2, sion. The diffraction peaks of ZrB2 were shifted to the right of the expected positions, indicating lattice contraction. The pres indicating lattice  expan ence of the shoulder corresponding to the position of the stoichiometric ZrB2 at 2y 5 41.61, on the left of the shifted ZrB2 peak on the pattern of the material containing 20 and 30 vol% TaSi2, probably indicates insufficient processing time to homogenize  the  phase  composition. The calculated unit cell volume from 30.70 to 30.02 A˚ 3 and increased for  decreased for ZrB2  1100  1200 1300 Temperature (°C)  1400  0  40  80  120  160  200  ZrB2  8 v.% Ta5Si3  15 v.% Ta5Si3  31 v.% Ta5Si3  25 v.% SiC  Scale Thickness 115 µm 1 300 µm  2  2  1  M  a  s s  C  e g n a h  (  g  /  m  2  )  Fig. 6. Mass changes during furnace oxidation (2 h) of ZrB2 ceramics containing 0, 8, 15, and 31 vol% Ta5Si3 at 11001-14001C.  Fig. 7.  SEM/EDS of the surface of ZrB2/8 vol% Ta5Si3 ceramics after oxidation at 13001C for 2 h.  2254  Journal of the American Ceramic Society—Talmy et al.  Vol. 91, No. 7      \\x0c', 'July 2008  High-Temperature Chemistry and Oxidation of ZrB2-Based Ceramics  2255  ceramics  increased  The oxidation resistance of ZrB2-TaSi2 with increasing TaSi2 content (Fig. 10). Samples containing 10 vol% TaSi2 showed higher mass gain than pure ZrB2 at all temperatures. This behavior was also observed for the ZrB2/5 vol% Si3N4 ceramics and was due to the significant evaporation of B2O3 from pure ZrB2 above 12001C, decreasing the total weight gain from the formation of ZrO2 and B2O3 during oxidation. The ceramics containing 20 and 30 vol% TaSi2 exhibit significantly less mass gain than ZrB2 up to 14001C. The matching values of the mass gain for ZrB2 and ceramics with 30 vol% TaSi2 at 15001C are again the result of extensive evaporation of B2O3 from ZrB2 at this temperature. This is supported by the thickness of the oxidized layer for ceramics containing 30 vol% TaSi2 after oxidation at 14001 and 15001C (50 and 60 mm compared with 300 and 600 mm for pure ZrB2, respectively). A small increase in the thickness of the oxidized layer for the TaSi2-containing material between 14001 and 15001C together with 5-10  times higher thickness for pure ZrB2 at additional and conclusive indications of the protective capabil these temperatures are  ities of Ta2O5-containing borosilicate glass. The SEM/EDS of the oxidized surfaces of all the ZrB2-TaSi2 samples showed glass immiscibility (similar to that for the ZrB2-Ta5Si3 ceramics) with a Zr-containing glassy phase and the crystalline phase contain ing mostly Zr with a small amount of Ta. The XRD of  the  oxidized surface of  showed the presence of ZrO2 and TaZr2.75O8. No Siand B-containing crystalline phases were identified.  the samples  Figure 11 shows the behavior of the most oxidation-resistant  compositions from each series of Si-containing additives char acterized in this research together with oxidation data previously reported for ZrB2/25 vol% SiC modified with 10 mol% TaB2.16 An inset presents a magnified picture corresponding to the mass gains at 12001-14001C. All the compositions exhibited improved oxidation resistance at 12001 and 13001C compared with the  ZrB2-SiC ceramics. The improved oxidation resistance persists with some of the compositions approaching 14001C, with the  best performance  shown by the ZrB2/SiC/TaB2 material. As discussed above, the lower oxidation resistance of the SiC-containing ceramics below 13001-14001C compared with all other  materials can be attributed to the insufficient oxidation of SiC at  these temperatures to form a continuous protective borosilicate  glass, while the significant oxidation of ZrSi2, Ta5Si3, and TaSi2 occurred at much lower temperatures. The presence of transi tion-metal oxides in the borosilicate glass,  inducing immiscibil Fig. 8. X-ray diffraction of ZrB2/10, 20, and 30 vol% TaSi2 ceramics HP at 20001C for 20 min.  TaSi2 from 130.19 to 132.67 A˚ 3 in the ceramics containing 20 vol% (15 mol%) TaSi2, confirming the formation of solid solutions. The solid solubility in the system was not expected and was  not previously reported in the literature. The crystal structures  of both ZrB2 and TaSi2 are hexagonal but belong to different space groups (ZrB2 is AlB2-type (P6/mmm) and TaSi2 is CrSi2type (P6222)). Based on the results of Brukl’s research,34 which excluded solubility of silicon or silicides in ZrB2, one way to solid solutions between ZrB2 and TaSi2 is to consider that only the transition metals mutually dissolve, leaving the boride and silicide lattices intact. With atomic  formation of  interpret  the  radii of Zr and Ta being 0.160 and 0.146 nm, respectively,  the  incorporation of Ta into the ZrB2 crystal TaSi2 crystal lattice had to result in the observed contraction of the ZrB2 and expansion of the TaSi2 crystal A mutual solubility of ZrB2 and TaSi2 is supported by the results of SEM/EDS studies, which showed that the majority of  lattice and Zr into the  lattices.  the crystals had a core-shell structure with the core being ZrB2 and the shell containing both Ta and Zr (illustrated on Fig. 9)  for ceramics with 30 vol% TaSi2. simultaneous presence (Fig. 8) of the shifted ZrB2 peaks and the stoichiometric ZrB2 peak reflects the core/shell structure of the material. It is expected that an increase in holding time during  It can be assumed that  the  hot pressing would complete the formation of the solid solutions  ity,  additionally  increased  glass-protective  capability.  The  and eliminate the core-shell structure of the crystals.  diminishing  effect  of  these  oxides  at  temperatures  above  Fig. 9.  SEM/EDS of ZrB2/30 vol% TaSi2 ceramics HP at 20001C for 20 min, BS image.  \\x0c', '14001C and the increase in oxidation rates could be related to a  decrease  in glass  viscosity, which leads  to the  reduction of  immiscibility  effects  and  an  increase  in  oxygen  diffusivity  through the glass layer.  IV.  Summary  Silicon carbide is conventionally used to improve the mechanical  properties and oxidation resistance of ZrB2 ceramics. The effect of SiC on the oxidation resistance of ZrB2 is due to the formation of protective surface borosilicate glass, which enhances  oxidation protection by acting as a barrier to the inward diffu sion of oxygen. The effect of Si3N4, Ta5Si3, and TaSi2 (as sources of Si instead of SiC) on the phase composition,  microstructure, and oxidation behavior of ZrB2 ceramics was characterized. Silicon nitride decomposed during the processing  of ZrB2-Si3N4 compositions, and the final ceramics contained ZrB2, BN, ZrSi2, and ZrN. Oxidation resistance of Si3N4-modified ceramics increased with increasing Si3N4 content and was further improved by the addition of Cr and Ta diborides due to  phase separation of the surface glass. The ZrB2/Ta5Si3 materials exhibited complex chemical interactions and solid solubility  between the  components during hot pressing. Limited solid  solubility was also observed between ZrB2 and TaSi2. Both Ta silicides significantly improved oxidation resistance of pure ZrB2 ceramics, with the best performance exhibited by ZrB2 containing 15 vol% Ta5Si3 and 30 vol% TaSi2. The effect is attributed to the pronounced immiscibility in the surface glass formed  during the oxidation tests. All studied compositions exhibited improved oxidation resistance below 14001C compared with 15001C,  ZrB2/SiC. At  the  SiC-modified  ZrB2  composition  performed  best,  but  was  only  slightly  better  than  the  ZrB2/SiC/TaB2 and ZrB2/Si3N4/CrB2 ceramics.  Acknowledgments  The authors would like to acknowledge ONR (Drs. S. Fishman and D. Shifler)  for providing support for this project and Dr. B. Varshal for valuable discussions  on phase separation in glasses.  References  1M. M. Opeka, I. G. Talmy, and J. A. 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Funayama, and T. Isoda,  ‘‘Reaction between Newly-Developed Si3N4 Fiber Mater. Sci. Lett., 19 [13] 1179-83 (2000). 25P. Shaffer,  and Various Ceramics,’’  J.  ‘‘An Oxidation Resistant Boride Composition,’’ Ceram. Bull., 41  [2] 96-9 (1962). 26P. Shaffer,  ‘‘Oxidation-Resistant Ceramics and Method of Manufacturing  Same’’; U.S. Patent 3, 189, 477 (1965) 27V. A. Lavrenko, A. D. Panasyuk, T. G. Protsenko, V. P. Dyatel, E. S.  Lugovskaya, and E. I. Egorova, ‘‘High-Temperature Reactions of Materials in the  ZrB2-ZrSi2 System with Oxygen,’’ Sov. Powder Metal Metal Ceram., 21 [6] 471-3 (1982). 28A. D. Panasyuk, L. A. Glebov, T. G. Protsenko, K. M. Timofeeva, and V. P.  Dyatel,  ‘‘High-Temperature Oxidation of Materials in the System ZrB2-CrSi2’’; pp. 127-30 in Vysokotemp. Boridy Silitsidy, Edited by T. Ya. Kosolapova. Nayk ova Dumka, Kiev, USSR, 1978. 29H. Pastor and R. Meyer, ‘‘An Investigation of the Effect of Additions of Metal  Silicides on Titanium and Zirconium Borides  from the Point of View of  their  1200  1300 1400 Temperature (° C)  1500  0  100  200  300  400  500  M  a  s s  C  e g n a h  (  g  /  m  2  )  ZrB2  10 v% TaSi2  20 v% TaSi2  30 v% TaSi2  ZrB2 / 25 v% SiC  1  2  50 µm 300 µm 60 µm 600 µm   Scale Thickness  2  1  3  4  3  4  Fig. 10. Mass changes during furnace oxidation (2 h) of ZrB2 ceramics containing 10, 20, and 30 vol% TaSi2 at 12001-15001C.  1200  1300  1400  1500  Temperature (° C)  0  100  200  300  M  a  s s  C  e g n a h  (  g  /  m  2  )  1400  1300  1200  20  40  60  ZrB  25 v % SiC  15 v % Ta Si  35 v % Si N  30 v % TaSi  20 v % Si N / 10 mole % CrB  25 v % SiC/ 10 mole % TaB  Fig. 11. Mass changes during furnace oxidation (2 h) of ZrB2 ceramics containing SiC, Si3N4, TaSi2, and Ta5Si3 at 12001-15001C.  2256  Journal of the American Ceramic Society—Talmy et al.  Vol. 91, No. 7          \\x0c', 'July 2008  High-Temperature Chemistry and Oxidation of ZrB2-Based Ceramics  2257  Sintering Behavior and their Resistance to Oxidation at High Temperatures,’’ Rev.  Int. Hautes Temp. Refract., 11 [1] 41-54 (1974). 30E. Opila, S. Levine, and J. Lorincz, ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39 [19]  5969-77 (2004). 31I. Talmy, J. Zaykoski, M. Opeka, and A. Smith,  34C. E. Brukl,  ‘‘Ternary Phase Equilibrium in Transition Metal-Boron-Car bon-Silicon Systems. Part II. Ternary Systems, Volume X’’; AFML-TR-65-2, Part  II, Vol. X, Air Force Materials Laboratory, Wright Patterson Air Force Base,  Dayton, OH, November 1966. 35R. Kieffer and F. Benesovsky, ‘‘Recent Developments in the Field of Silicides  the System ZrB2-Ta5Si3,’’ J. Mater. Res., 21 [10] 2593-9 (2006). 32T. J. Rocket and W. R. Foster, ‘‘Phase Relations in the System Boron Oxide-  Silica,’’ J. Am. Ceram. 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},{
  "_id": 88,
  "PDF": "High-Temperature Isothermal Oxidation of Ultra-High Temperature Ceramics Using Thermal Gravimetric Analysis.pdf",
  "Text": "['High-Temperature Isothermal Oxidation of Ultra-High Temperature Ceramics  Using Thermal Gravimetric Analysis  Melia Miller-Oana and Erica L. Corral  †  Materials Science and Engineering Department, Arizona Materials Laboratory, The University of Arizona, 1235 East James E  Rogers Way, Tucson, Arizona 85721  Oxidation  of  ZrB2 + SiC isothermal measurements to study the  composites  is  investigated  using  effects of  temperature,  time, and gas flow on oxidation behavior and microstructural  evolution. A test method called dynamic nonequilibrium ther mal gravimetric analysis  (DNE-TGA), which eliminates oxida tion during the heating ramp, has been developed to monitor  mass  change  from the  onset  of  an isothermal hold period (1000°C-1600°C) and gas  (15 min) as a function temperature  flow (50  and  200 mL/min).  In  comparing  isothermal  to  non isothermal  TGA measurements,  the  scale  thicknesses  from  isothermal  tests are up to 4 times greater,  indicating that oxi dation kinetics are faster for isothermal testing, where the oxide scale thickness is 110 lm after 15 min at 1600°C in air. Isothermal oxidation followed parabolic kinetics with a mass temperature dependent from 1000°C-1600°C. The mass gain increased from ~5 to 45 g/m2 and parabolic rate constants increased from 0.037 to 2.2 g2/m4\\x01s over this temperature range. The effect of flow velocity on oxidation is not sig gain that  is  nificant  under  the  given  laminar flow environment where  the  gas boundary layer is calculated to be 4 mm. These values are  consistent with diffusion of oxygen through the glass-ceramic  surface layer as rate limiting.  I.  Introduction  T HE  use  of many materials  is  restricted  by  undesirable  reactions  in  high-temperature  oxidizing  environments.  Ultra-high temperature  ceramics  (UHTCs) are high melting  temperature materials that include transition-metal borides, carbides, and nitrides1 that, because of their resistance to oxi dation  at  high  temperatures,  are  attractive  candidates  for  applications  such  as  jet  engines,  gas  turbine  engines,  and  thermal protection system materials  for hypersonic vehicles.  This oxidation resistance  is due  to protective  surface  layers  that form initially at temperature and retard further penetration of oxygen.2-5 Methods representative use conditions  to study oxidation of UHTCs  in  are  limited, which makes  it  challenging to develop an understanding of  the  response of  UHTC materials  to  different,  interacting,  high-temperature  process variables. This article reports use of a readily avail able thermogravimetric analyzer (TGA) temperature: 1600°C) to investigate oxidation kinetics ZrB2 + SiC composites as a function of different hold temperatures.  (maximum operating  for  isothermal  Thermogravimetric analyzer  furnaces are  commonly used  for oxidation testing where mass gain is measured as a functemperature in an oxidizing environment.6-10 Typical TGA heating rates (5°C-20°C/min) can allow significant  tion of  UHTC surface reactions on the heating and cooling shoulders of any isothermal holds.4,7,9,11-14 We furnace-based oxidation test method as conventional  label  this general  testing.  Two ways  to perform isothermal oxidations without  the  heating  and  cooling  artifacts  are  (1)  quickly  inserting  the  specimen at  the desired temperature  and (2) heating  in an  inert environment and then switching to the reactive gas once  the desired temperature is reached. Obtaining mass change as  a  function of  time  at  temperature with method (1) usually  requires periodic removal of  the specimen, cooling it, weigh ing, and then reheating to the desired test temperature. Shugart et al.11 used the first method to isothermally oxidize ZrB2 + 30 vol% SiC in a box furnace from 1300°C-1550°C for up to 100 h to determine the weight change and the oxide composition. Levine et al.4 studied the dation behavior of ZrB2 + 20 vol% SiC 1327°C and 1627°C for 100 min. However, reported in  scale  isothermal oxi in  furnaces  at  it is also imporsitu mass  tant  to note  that neither  study  gain;  thus,  leaving uncertainty about oxidation kinetics.  One form of method (2) is to use a TGA furnace that allows  the specimen to be heated in an inert gas environment that can  be quickly switched to air once the isothermal test temperature is reached. Chen et al.15 oxidized hot-pressed LaB6-ZrB2 from 912°C-1223°C, using a TGA furnace with a heating rate of 100°C/min in argon for up to 80 min in order to determine the activation energy of oxidation. Guo et al.16 oxidized ZrB2 powder from 650°C-800°C in a TGA furnace for up to 8 min to determine the kinetic rate constants. Rioult et al. used a  similar experimental technique to study the isothermal oxidation of Mo-Si-B specimens in a TGA furnace.17  Previously, we reported experiments on UHTCs in a TGA  furnace with oxygen present only at  the isothermal set temchanges.18  perature  and  semicontinuous  recording of mass carbon-carbon  Those  studies  on UHTC-filled  composites  showed that because the carbon matrix did not oxidize in the  high-purity Ar  atmosphere  during  the  heating  ramp, we  could focus on the behavior.18 As has been reported previously,19 oxidation of ZrB2 and SiC are thermodynamically favored at all temperatures in air, but reaction rates are not appreciable below 700°C for and 1200°C for SiC. Possible oxidation products ZrB2 ZrO2, B2O3, SiO2, SiO, CO2, and CO. At high temperature, SiO2 and B2O3 are viscous liquids that form adherent surface coatings that protect the underlying composite  effect of filler  composition on oxidation  are  from further  oxidation. The vapor pressure of B2O3 becomes above 1050°C and volatile SiO is favored over SiO2 at partial pressures of oxygen, both of which properties affect  significant  low  oxygen transport  through the surface layer. Reaction kinetics  are  important  for  these  short-time  experiments because the 800°  ZrB2 oxidation rate C-1200°C. Thus,  is  faster  than  that  of  SiC at  the  difference  in  oxidation  kinetics will  influence the composition of at T > 1200°C) where  the protective oxide  (a viscous  liquid  SiO2 diffusion  is more  protective  than  B2O3 will be rate limiting.  13,18,20  and  oxygen  through  this  liquid  layer  B. Fahrenholtz—contributing editor  Manuscript No. 36590. Received March 25, 2015; approved September 22, 2015.  †  Author  to whom correspondence  should be addressed.  e-mail:  elcorral@email.ari  zona.edu  619  J. Am. Ceram. Soc., 99 [2] 619-626 (2016)  DOI: 10.1111/jace.14001  © 2015 The American Ceramic Society  Journal  \\x0c', 'The stoichiometric reactions are:  ZrB2 þ 5=2O2 (g) ! ZrO2 þ B2O3 (l)  (1)  SiC þ 3=2O2 (g) ! SiO2 (g) þ CO(g)  (2)  B2O3 (l) ! B2O3 (g)  (3)  SiC þ O2 (g) ! SiO(g) þ CO(g)  (4)  Here, we report mass changes ZrB2 + 25 vol% SiC specimens for short times high temperatures (1000°C-1600°C), and relate high-temperature gas flow velocity profiles  on  oxidizing monolithic (≤15 min) at  them to cal culated  around  the solid specimen and to the observed changes  in specimen  microstructure  and  composition  over  time. We  tested  for  short  times because we are specifically interested in how fast  these oxides form in the early stages of UHTC oxidation and  in the  complex  layered oxide  structures  that  are produced.  These  results  are directly  compared to those  from conven tional TGA measurements  to provide  further understanding  of  the formation and growth of  the layered oxide scale. We  compare directly to the conventional TGA measurements  so  that we have a framework that has been studied as a function of temperature. As reported by Rezaie et al.7 during conventional TGA oxidation of ZrB2 + 30 vol% SiC up to 1500°C in air, mass gain began below ~800°C due to ZrB2 oxidation where ZrO2 and B2O3 formation and growth occurred. Between 1215°C and 1300°C, mass loss is observed  and is attributed to B2O3 SiC began oxidization to form SiO2. At even higher temperatures mass gain is observed again, which is attributed to the  evaporation, while  simultaneously  continued formation of SiO2, which is occurring faster B2O3 evaporation; however, B2O3 volatilization continues resulting in a more SiO2-rich oxide scale; which will oxygen diffusion more than a B2O3-rich oxide scale.20  than  limit  II.  Experimental Procedure  (1)  UHTC Preparation  Powder processing and ZrB2 + 25 vol% SiC by Walker et al.21 and are ZrB2 + 25 vol% SiC powder hexane, rotoevaporated and 1800°C for 20 mm inner diameter graphite dies. Final dense ~2 mm thick. After  spark  plasma  sintering  of  dense  specimens  are  previously  described  briefly  summarized  here.  is ball-milled with WC balls  in  sieved  before  spark  plasma  sintering  at  5 min  under  35 MPa  using  12  or  specimens  for  both  diameters  are  densification,  specimens are ground to remove adhering graphite foil using  resin-bonded 220 grit diamond pads. Densities are determined by the Archimedes method (all specimens ~99% theo retical  density).  Before  oxidation  testing,  specimens  are  sectioned  into  bars or remained cylinders, 125 lm and washed acetone and then with distilled water. The geometric dimen polished  to  a  surface  finish  of  ultrasonically  using  sions and mass (0.1 mg accuracy) are recorded before oxida tion testing.  (2)  Oxidation Testing Using Dynamic Nonequilibrium  TGA (DNE-TGA)  Cylindrical specimens are ~12 mm diameter and 2 mm thick, and one bar specimen (tested only at 1600°C) has dimensions of ~2 mm 9 7 mm 9 9 mm. A given test specimen is placed  directly on a ZrO2 disk as shown in Fig. 1(a) inside the TGA furnace (Netzsch STA 449F3 Jupiter, Selb/Bavaria, Germany). Specimens are tested at 1000°C-1600°C in air  for up  to 15 min. The  free upper  surface of  the  specimen and the  lower  one  against  the ZrO2 local environments in the upward flowing gas.  support  experienced  different  The  experimental procedure is similar and Corral.18  to one previously  described by Walker (20°C/min)  Specimens  are  heated  in  ultra-high  purity  argon  (UHP Ar)  to  the  isothermal  temperature with  a  flow rate  of  100 mL/min.  Once  the  desired  temperature  is  reached,  the  system is  allowed to equilibrate for 2 min before the gas is switched to  air  (200 mL/min). After holding for (1000°C-1600°C), back to UHP Ar with a flow rate of 50 mL/min and allowed  up  to  15 min  at  the  desired  temperature  the  gas  is  switched  to cool freely. The purity of the UHP Ar is 99.999% with ~1 ppm by volume of oxygen. The gas environment changed  from inert  to oxidizing  in less  than 4 s. The mass balance  data  acquisition  rate  is  600 points/min  during  the  experi ment, which gave a semicontinuous record of mass change the mass balance is 1 lg.  with time. The digital resolution of  Due  to the  low level oxygen impurity  in UHP argon,  a  SiC-depleted region is  formed on the UHTC outer  surface  due to active oxidation 1300°C. The effect of the outer SiC-depleted region on mass gain at 1200°C is determined by comparing two specimens: first one is heated in Ar at 1300°C to form the depleted layer  of  SiC (as  SiO gas)  beginning  at  and the other without  the preheating that  lacked the depleted  layer. The two specimens are isothermal TGA tests at 1200°C.  then  subjected  to  identical  (3)  Oxidation Testing Using Conventional TGA  A ZrB2 + SiC bar (~2 mm 9 5 mm 9 11 mm) specimen is heated at 20°C/min in air to 1600°C with an inlet gas flow  rate of 50 mL/min and held at  that  temperature  for 15 min  in the same TGA furnace used for DNE-TGA.  In this case,  Fig. 1.  (a) Experimental  setup  for  the TGA used  in  isothermal  oxidation. Velocity flow maps  from (b) DNE-TGA method using a  gas flow of 200 mL/min.  620  Journal of the American Ceramic Society—Miller-Oana and Corral  Vol. 99, No. 2  \\x0c', 'the specimen rested on an additional ZrO2 setter as shown in Fig. S1. Testing at 1600°C is selected to maximize any effects  for  comparison to the DNE-TGA results. After  the isother mal hold, specimens are allowed to cool  freely to room tem perature in the presence of air.  To compare  the  conventional  and DNE test methods,  a  third test  is performed where a combination of both methods  is used, first to heat a specimen in air using the conventional method to 1600°C in order 500°C. After switching to flowing Ar gas, the DNE-TGA test is then done by the usual method at 1600°C  to pre-oxidize  it, and then it  is  cooled  to  for a 15 min hold.  (4)  Multiphysics Modeling of  the Gas Flow Environment  in  the TGA Furnace  A multiphysics  software  program (COSMSOL, the Navier-Stokes  Inc.,  Los  Angeles, CA)  that  solves  equation  to  model  the gas flow inside the TGA furnace is used to deter mine the boundary flow conditions around the specimen dur ing  testing.  The model (1600°C),  assumes  that  the  temperature  is  isothermal  the  pressure  is  atmospheric,  and  the  flow is  laminar.  In the TGA furnace,  the gas  is  introduced  through the bottom inlet with a flow of 0.2 or 0.8 m/s (50 or  200 mL/min) and exited the outlet at  the top [Fig. 1(a)]. The  maximum inlet  flow velocity  using  the  current  equipment  configuration is 200 mL/min. High-temperature gas velocity  maps are constructed to compare the local gas flow velocity  and boundary  layer  thickness  around the  specimen  during  conventional and DNE-TGA testing using 50 and 200 mL/  min inlet flow velocities, respectively.  (5)  Microstructural Characterization  After  oxidation  testing,  the  specimens  are  polished  in DI  water on SiC pads down to a mirror finish so that  the thick ness of  the oxide  scale on cross  sections of  the  specimens  could be measured and the oxide  structure  could be  exam ined using scanning electron microscopy (SEM; S-4800, Hita chi  Inc., Pleasanton, CA) and energy-dispersive spectroscopy  (EDS; ThermoNORAN NSS, West Palm Beach, FL). The  polishing procedure  consisted of mounting cross  sections  in  epoxy and polishing successively on 80 grit, 120 grit, 220 grit, 40 lm, and 25 lm diamond polishing pads before  the final  polish with a 1200 grit SiC pad. Specimens are imaged in sec ondary electron mode using both the upper and lower detec10-20 kV  tors  with  accelerating  voltages  of  at  various  magnifications.  Using  ImageJ  software  (NIH,  Bethesda,  MD), oxide scale thicknesses are measured on one image of  each condition where the thickest oxide layer  is present. The  oxide layer is thicker on the top surface of  the specimen than  the bottom due to more gas flow interaction, as  is predicted  by our gas flow velocity models.  In addition,  ImageJ is used  to determine grain size by counting at  least 300 grains  for a  representative specimen.  III.  Results and Discussion  (1)  Gas Flow Environment and Reliability of Mass  Measurements  The  calculated gas velocity maps  in the TGA furnace using  inlet  gas  velocities  of  0.2  and  0.8 m/s  are  compared  in  Fig. S1  and 1(b),  respectively. They  show that because of  boundary layer effects,  the gas velocities at  the specimens are (<0.001 m/s). The result in similar lami the  same  for  both TGA configurations  velocity maps  for both inlet velocities  nar gas flow where the gas flow is  slowest around the speci men. Although the  inlet gas velocities differ by a factor of  four,  there is no significant difference in gas flow velocity at  the  specimen because the ~4 mm, which  gas  boundary  layer  around  the  specimen  is  is  about  1/6  of  the  inner  tube  diameter of  the furnace. Therefore,  the flow environment at  the  specimen and chemical  reactions  at  the  surface  can be  treated the same for either DNE-TGA or conventional TGA.  The reproducibility in the mass measurements is shown in Fig. 2(a) where 3 specimens are oxidized at 1600°C using  the DNE method. The mass gains after 8 min in the isothermal hold are 1.00% (37.2 g/m2), 1.11% (33.8 g/m2), 0.96% (35.2 g/m2); 35.4 \\x06 1.70 g/m2. The normalized mass gains do not therefore, the average mass gained with mass percentage gains due to the differences in surface  and  is  trend  areas. The parabolic rate constants determined by linearly fitting the square of mass changes (g2/m4) as a function of are 2.3345, 2.3805, and 2.1783 g2/m4\\x01s, giving an average of time 2.30 \\x06 0.11 g2/m4\\x01s, which shows that the reproducible data are obtained using DNE-TGA. However,  there is an experi mental variability associated with testing specimens of differ ent  surface areas  in conventional TGA as shown in Fig. S2, tested at 1600°C for 15 min. The final mass gain for the smaller specimen (48.5 mm2) is 1.1% (30.9 g/m2), whereas for the larger specimen (171 mm2) mass gain is 0.40% (14.1 g/m2). We hypothesize that because smal where  specimens are  ler  specimens have higher edge length to surface area ratios;  more mass gain is gained due to more oxidation occurring at  the  edges and corners  than for  larger  specimens. Therefore,  to minimize this variable, cylindrical specimen surface areas are kept between 272 and 307 mm2, and bar face area is 171-195 mm2.  specimen sur The uncertainty is determined for  the mass measurements  obtained  in  Fig. 2(a)  where  the  mass  gain  average  is  Fig. 2.  (a) Mass as a function of  time using DNE-TGA at 1600°C  showing  consistency  in measurements.  (b) Mass measurements  for  conventional heating, 1600°C  combination of DNE and conventional,  and  DNE  at  showing  that  when  a  conventionally  heated  specimen is tested using DNE-TGA,  it will not exhibit as much mass  gain as  the DNE method due  to the  formation of  the oxide  scale  formed during conventional heating.  February 2016  Isothermal Oxidation of UHTCs  621  \\x0c', '2.73 \\x06 0.75 g. The uncertainty is calculated from the resolution of the mass balance, the standard deviation of the 3  measurements, and mass  signal precision of 0.5%. The aver2.73 \\x06 0.87 g with  age with an expanded uncertainty 95% level of confidence (k = 2).  is  a  Figure 2(b)  shows  the mass measurements when a speci men is first oxidized using  the  conventional method (solid  line) and then switched to the DNE method (long dashes).  That specimen, which is labeled as “Combo” is compared to  a specimen oxidized using only the DNE method (short dashes) at 1600°C. The Combo and DNE curves have been  moved to the end of the conventional heating for comparison  purposes, meaning that Combo and DNE curves start where  the Conv  curve  ends. As  expected,  the  specimen  oxidized  using the combination method behaves  like the conventional  TGA method and does not exhibit as much mass gain as the  specimen  oxidized  using  the DNE method  because  of  the  oxide layer formed during conventional heating that acts as a  protective barrier. the combination method is 0.775 g2/m4\\x01s, which is In addition, the parabolic rate constant for than that of the DNE-TGA method (2.22 g2/m4\\x01s).  lower  (2)  Conventional TGA and DNE-TGA Mass  Measurements and Microstructures  Figure 3(a)  shows mass measurements as a function of time 1600°C for  for  conventional TGA and DNE-TGA held at  15 min. The specimen oxidized in air using the conventional  method is  consistent with mass  change  trends discussed in et al.7 Reaction  the  introduction such as those for ZrB2 + 25 SiC oxidation have shifted temperatures as compared to ZrB2 + 30 SiC, but same mass gain and loss regimes are observed. Oxygen diffu of Rezaie  temperatures  to  lower  the  sion will be  slower  in a SiO2-rich outer in less mass gain using conventional TGA compared  layer, which should  result  to DNE-TGA. Figure 3(a) shows that less mass is gained during conventional TGA (~0.4% or 14.1 g/m2) as compared to DNE-TGA (~1% or 45 g/m2) 1600°C.  after  the  15 min hold at  Figure 3(a) also shows  the mass change versus  time using  the DNE-TGA method. The specimen oxidized using DNE TGA did not begin to gain mass until 1600°C is reached. However, mass that, between ~1300°C-1600°C, due  the  isothermal  test  temperature  of  loss  occurred before  to the  active oxidation of SiC [Reaction (4)]  that  is  favored in the  low partial pressure of oxygen present  in UHP Ar gas. This  SiC-depleted region is an artifact of  the  test method and it  did not significantly influence oxidation measurements, shown in Fig. 3(b) where at 1200°C, tion of time is not dependent on  as  the mass gain as a func an  outer  SiC-depleted  region.  In particular, the specimen with a SiC-depleted layer 11.5 g/m2, whereas the specimen without 9.50 g/m2. While  gained  the  SiC depleted  layer  gained  there  is  a  21%  increase from the specimen without  the SiC-depleted layer to  the one with a SiC-depleted layer, the difference is only 2.0 g/ m2 and similar to the standard deviation (1.70 g/m2) calculated at 1600°C after 8 min for the TGA reliability measure ments. When the specimen without  the SiC-depleted layer is the solid-  exposed to air, oxygen will  react with the solid at  gas  interface. However, when the specimen with the SiC-de pleted layer is exposed to air, solid-gas  in addition to oxygen reacting  at  the  interface,  oxygen may  be  able  to  diffuse  through  the  SiC-depleted  region more  quickly.  This will  depend  on  oxygen  pathways  formed  in  the  SiC-depleted  region; the connectivity of the grain-boundary network is dis cussed later. We assume  that  the  effect of an outer SiC-de pleted layer on the UHTC composites will not be significant temperatures >1200°C where  at  the  same mix of oxidation  and vaporization reactions occurs  in the  entire  temperature  range. However,  future  experiments  should be designed to  determine  the  effect of  inert heating at higher  temperatures.  As the reaction rates increase at higher temperatures, oxygen  diffuses more rapidly through the outer SiC-depleted region in the composites (25 lm thick at 1600°C) and reactions (1)-  (3) occur simultaneously and rapidly, resulting in the formation of an outer ZrO2-SiO2-B2O3 liquid. By studying the temperature dependence of mass changes in DNE-TGA and the 1000°C-1600°C, we  resulting microstructures  at  can  deter Fig. 3. (a) Conventional TGA and DNE-TGA in situ mass measurements. (b) Effect of SiC-depleted layer image of conventional TGA oxidation and (d) SEM image of DNE-TGA oxidation at 1600°C in air  in DNE-TGA at 1200°C.  (c) SEM  for 15 min. The outer oxide layer consists  of SiO2-based glass and the intermediate layer consists mostly of ZrO2 with some glass.  622  Journal of the American Ceramic Society—Miller-Oana and Corral  Vol. 99, No. 2  \\x0c', 'mine when SiC oxidation becomes kinetically favorable, and  how the scale forms and grows over this temperature range as  a function of time.  In comparing the conventional and DNE-TGA mass mea surements  shown in Fig. 3(a), we can see that (~700°C) tional method as compared to the DNE-TGA method, which oxygen is not admitted until 1600°C.  the mass gain  began  at  a  lower  temperature  using  the  conven in  Instead UHTCs  lost mass during initial heating due to active oxidation of SiC.22  in DNE-TGA in UHP Ar  In conventional TGA dur ing the heating period, mass evaporation,7 whereas mass  loss  is observed due  to B2O3 in DNE loss  is  not  observed  TGA once  air  is  introduced into the  furnace. Because  the  conventional TGA specimens are heated entirely in air,  the  oxide layer  is able to form during this  initial heating period,  and the oxide  scale  formed during DNE-TGA will differ  in  layered composition from one formed during a conventional  TGA test.  Figure 3(c)  shows an SEM image of  the oxide scale on a  cross section of the specimen oxidized using the conventional method at 1600°C for 15 min in air. EDS maps of O, Zr, and Si for the same specimen are shown in Figs. S3(a)-(c).  Note  that  boron  and  carbon  cannot  be  reliably  detected  with  this  SEM/EDS method,  so  any  statements  on  boron  and  carbon  compositions  are  inferences  based on 1600°C,  indirect  evidence. After  holding  for  15 min  at  the  outer  oxide scale consists mostly ~6 lm thick. The the bulk consists of Zr and O (as ZrO2) without Si is ~30 lm thick as shown by EDS. Figure 3(d) shows  of  Si  and O (as  SiO2) layer between the outer oxide scale  and  is  and  (as SiC)  an  SEM image  of  the  oxide  scale  on  a  cross  section  of the 1600°C are pre specimen oxidized using  the DNE-TGA method at  for  15 min  in  air. EDS maps S4(a)-(c). The  for O, Zr,  and  Si  sented  in Figs.  images  show that  after  the  15 min hold,  the outer oxide  scale, which consists of Si, O, ~25 lm thick. (SiO2 and ZrO2 Although not detected, B is likely present and we infer that  and  some Zr  particles)  is  its  concentration layer.14 The  changes  with  depth  through the outer ~35 lm thick  oxide  intermediate  layer  is  and  consists of Zr, O, and some Si  (mostly as ZrO2 with small ~50 lm SiC-depleted region is  amounts  of  SiO2); consists mostly of ZrB2. Carney Si-O-C-B phase using transmission electron microscopy after oxidation of ZrB2 + 20 SiC at high tem(1400°C-1600°C) for 150 min in air.12 This result  and  the  thick  and  et al.  reported  observing  a  peratures  supports  the idea that not all C and B are lost due to reac tions  and evaporation and that B concentration is  a  func tion of oxide depth.  In comparing the oxide  layers  formed on specimens oxi dized using conventional and DNE-TGA, we observed that formed during DNE-TGA is ~4 times  the outer oxide  scale  thicker than the diffuses ~3 times  conventional TGA scale  and that oxygen  farther  into the bulk ceramic during DNE TGA. Both methods  result  in  the  formation  of  a  layered  oxide  scale; however,  the  thickness and composition of  the  layers differ. The outer oxide scale (mostly SiO2) formed during conventional TGA is ~6 lm thick, and outer oxide scale consisting of a glass with ZrO2 particles formed during DNE-TGA is ~25 lm thick. The conventional intermediate layer is ~30 lm thick while the DNE-TGA intermediate layer is ~35 lm thick. No separate SiC-depleted region is during conventional TGA, whereas during DNE-TGA, SiC-depleted region is ~50 lm thick.  formed  the  (3)  DNE-TGA Oxidation and Parabolic Kinetics  Figure 4(a)  shows mass gains from DNE-TGA 1000°C-1600°C in isothermal periods. Mass increases with increasing time due to the simultaneous oxidation reactions [Reactions (1)-(3)] of  heating  experiments  at  air  during  the  15 min  ZrB2 and SiC, and mass increases with increasing temperafrom 1000°C-1600°C, between 1200°C-1300°C. ture except At 1300°C, slightly less mass is gained than at 1200°C, which  is  likely due  to the  formation of a liquid SiO2 with more B2O3 evaporation. Because the viscous SiO2 is a more protective barrier than B2O3, it limits oxygen diffusion and slows mass gain, resulting in less mass 1300°C 1200°C.23 1000°C-1300°C after the from ~5-10 g/m2, and from gain ranges from ~15-45 g/m2. For the mass gain (g/m2)  layer along  layer  gain  at  than  From  15 min hold, mass gain ranges 1400°C-1600°°C, mass  all  temperatures,  the slope of  is  steeper  at  the  beginning  of  the  hold  and  becomes  lower  as  time  increases, which  is  typical  for  parabolic  (diffusion  limited)  kinetics.  In addition,  the  change  in the  composition of  the  outer oxide  layer with time will  slow mass gain because as  more B2O3 protective.  evaporates,  the  resulting SiO2-rich layer  is more  The squared mass gain as a function of time from 1000°C-1400°C shown in Fig. 4(b) and from 1500°C-1600°C  in Fig. 4(c) fit a parabolic kinetic model. The squared mass gains are linear with time and have R2 values of >0.992 for from 1000°C-1600°C in DNE-TGA oxidation temperatures of ZrB2 + 25 SiC in air. Shugart et al.11 reported parabolic for oxidation of ZrB2 + 30 SiC from 1300°C- rate constants 1550°C for 100 min. They are in good agreement (same from 1300°C-1500°C with the parabolic  order of magnitude)  rate 0.05 g2/m4\\x01s constants reported here and the difference 1300°C-1400°C, observation of parabolic kinetics  is  less  than  at  as  shown  in Table I. The  and their  agreement with  published  values  is  strong  evidence  that  oxygen  diffusion  through the outer glass layer is the rate-limiting oxidation of ZrB2 + SiC using DNE-TGA.23  step  in  Fig. 4. (a) DNE-TGA in situ mass change measurements 1500°C-1600°C.  from 1000°C-1600°C. Square of mass change from (b) 1000°C-1400°C and from (c)  February 2016  Isothermal Oxidation of UHTCs  623  \\x0c', '(4) Microstructural and Compositional Analysis of DNETGA Specimens from 1000°C-1400°C  Figure 5 shows surface oxide layers on cross-sectioned speci1000°C, mens oxidized using the DNE-TGA method at 1200°C, and 1400°C. At 1000 and 1200°C, oxygen diffuses <25 lm into the supported by the outer reaction zone. The SiC grains are not  specimen, as  seen in Figs. 5(a) and (b) as  oxidized at  these  lower  temperatures  in the  isothermal hold  (or heating in argon);  therefore, no SiO2 glass sent. In Fig. 5(c), we observe an outer glass layer that resulted from oxidation of SiC at 1400°C. Figure 4(a), shows in mass gain from 1200°C to 1300°C, which we  layer  is pre a decrease  attribute to the temperature at which SiC begins  to oxidize.  Figure 5(c) supports this conclusion with evidence of SiC oxidation in the form of an outer glass (SiO2) layer at 1400°C.  (5) Influence of Time on DNE-TGA Microstructure Formation at 1600°C  Oxygen reactions  result  in a larger oxide scale for the DNE TGA specimen as  compared to the  conventional  specimen  [Figs. 3(c) and (d)];  therefore, we  conclude  that oxygen has  diffused farther  into the DNE-TGA specimen likely through reported by others.24-29 To determine scale grows and changes composition during  grain boundaries,  as  how the oxide  the isothermal period using DNE-TGA,  specimen cross  sec tions  are  examined  at  0,  8,  and  15 min  as  shown  in  Figs. 6(a), (b) and (c), respectively. The specimen oxidized for 0 min at 1600°C, which is exposed during heating only to  UHP Ar  (1 ppmv oxygen contaminant) has a SiC-depleted 25 lm thick. During  region  that  is  the  isothermal  period,  when  the  specimen  is  heated  in  air,  oxygen  can  diffuse  through the outer SiC-depleted region and react with ZrB2 and SiC to form oxide products. However, if the oxygen dif fuses along grain boundaries into the bulk material,  it will be (5),30  dependent on a percolation threshold. Using Eq. (V)  the  critical  percolation  volume  for  a  connected  SiC  microstructure can be calculated based on the grain sizes of  ZrB2 and SiC in the composite. V ¼ 100=ð1 þ :25ð/=Xc ÞðRz =Rs ÞÞ  (5)  where / is  the particle packing mode factor particle surface fraction (0.42), Rz is size (1.8 lm), Rs is size of SiC grains (0.80 lm).  (1.27), Xc of ZrB2  is  the  grains  From grain size measurements on the ZrB2 + SiC specimens, the critical volume for interconnected SiC is calculated therefore, ZrB2 + 25 vol% SiC is threshold and an interconnected SiC network will not composite.31 This means  to  be  37%;  below this  form  within the  that during heating  in  argon  in DNE-TGA,  oxygen must  diffuse  through  grain  boundaries  to  reach  SiC in  order  to  actively  oxidize  it.  Because oxygen does not have an interconnected network to  diffuse through the outer SiC-depleted region is limited to 25 lm in thickness after heating to 1600°C in argon. Immediately after oxygen is introduced, it will react with  ZrB2 form B2O3 ZrO2. Shortly after, oxygen will diffuse along grain boundaries and react with SiC to form SiO2 and CO.27,32 A borosilicate liquid with dissolved ZrO2 will protective barrier to limit further oxygen  in  the  outer  SiC-depleted  region  to  and  form and act as a diffusion.18  It  is  possible that  initially the liquid will not wet  the porous, SiC depleted outer layer. However, as discussed in Section III(2),  the mass gain on the  specimen with the SiC-depleted outer  layer  is not  significantly different  from one without  the SiC depleted layer. Therefore, we can conclude that the liquid formed at 1200°C on the SiC-depleted layer did not have dif ficulty wetting the specimen surface, and the liquid formed at  higher temperatures will easily wet ZrB2 and ZrO2 and act as a protective barrier.  For the DNE-TGA specimen (8 min) in Fig. 6(b), the outer SiO2 glass layer (~10 lm) contains some ZrO2 particles, whereas the intermediate layer is SiC depleted, contains ZrO2, and is ~10-20 lm thick. The EDS maps show that outer oxide scale is Si rich and the intermediate layer con the  tains mainly Zr and O, as shown in Fig. S5. During the last 7 min of DNE-TGA testing at 1600°C, specimens gained only ~10 g/m2. However, the oxygen penetration depth increased from ~30 to 110 lm. This may be  the  explained by the fact  that, even though the specimen is  still  gaining mass,  the rate (as determined by the change in slope)  has slowed due to increased evaporation of B2O3, which also leads to decreased mass gain due to slower oxygen diffusion  in  a more  SiO2-rich that slows diffusion because oxygen must  liquid  and  increased  thickness  of  the  oxide layers  travel  farther  to  reach  the  bulk material. The  outer  oxide  layer  formed  after  15 min  in DNE-TGA, which  contains ZrO2 the outer glass layer, is  particles across  the  entire  length of  shown in Fig. 6(c). ZrO2 particles are formed as B2O3 evaporates from the liquid and ZrO2 can no longer dissolve in the liquid resulting in the precipitation of the particles. The par ticles  are present  in the  glass  after  8 min, which  indicates  that  the  glass  layer  composition  continuously  loses B2O3, in the precipitation of ZrO2. The nucleation and growth of the oxide scale formed during DNE-TGA at 1600°C in air in Fig. 6(d). Oxygen must flow through the layer (~4 mm)  resulting  from 0 to 15 min is  summarized schematically  gas boundary  to the specimen in the TGA furnace. Initially,  a SiC-depleted region is present on the outer  surface. After  8 min,  the  layered oxide  structure  is  formed with an outer  glass  layer  containing ZrO2 is SiC-depleted and  particles  and  an  intermediate  layer  that  contains mostly ZrO2 with  Fig. 5. SEM images of cross-sectioned specimens oxidized using DNE-TGA with high magnification image as (c) 1400°C in air for 15 min. The images have the same scale.  inset at  (a) 1000°C,  (b) 1200°C,  Table I.  Parabolic Rate Constants  Temperature (°C)  DNE parabolic rate (g2/m4\\x01s)  constant  Shugart parabolic (g2/m4\\x01s)  rate constant  1000  0.037  — — —  1100  0.059  1200  0.095  1300  0.0852  0.11  1400  0.257  0.21  1500  0.885  0.12  1600  2.22  —  624  Journal of the American Ceramic Society—Miller-Oana and Corral  Vol. 99, No. 2  \\x0c', 'February 2016  Isothermal Oxidation of UHTCs  625  Fig. 6.  (a) SEM image of cross-sectioned specimen heated in UHP argon to 1600°C for 0 min resulting in SiC-depleted region as 1600°C for  8 min and (c)  EDS map. SEM images of cross-sectioned specimens oxidized using DNE-TGA at evolution of oxidation process at 1600°C from 0 to 15 min where initially at 0 min,  the outer layer is SiC-depleted. After 8 min,  15 min.  (b)  shown in Si  (d) Schematic of  the outer glass  layer,  intermediate  layer of glass and ZrO2, and the SiC-depleted regions have diffusion through the outer glass layer to the unoxidized material.  formed. After 15 min, all  layers grow resulting from oxygen  some  glass  oxides.  No  separate  SiC-depleted  region  is  oxidation testing are  compared to those obtained with con observed  after  8 min,  but  instead  the  outer  SiC-depleted  becomes the intermediate region. The outer oxide layer grows  due to formation of SiO2, B2O3, and ZrO2. After 15 min, layered oxide structure has continued to grow, resulting in a  the  thicker outer glass  layer of ZrO2 and glass, and a large SiC-depleted region that is mostly ZrB2.  layer, an intermediate  IV.  Summary  ZrB2 + 25 Composite SiC specimens are oxidized from 1000°C-1600°C using a thermogravimetric analysis technique  (DNE-TGA)  in which the specimens are exposed to flowing  air only during an isothermal hold period in order to obtain for ZrB2 + SiC oxidation temperatures using an isothermal method by eliminating the  constants  parabolic  high  rate  at  effect of oxidation during heating and isolate only the effect  of  isothermal period on mass measurements. The isothermal  mass measurements  and  oxide  scales  that  formed  during  ventional TGA performed with slow heating entirely in air where more mass is gained using DNE-TGA (45 g/m2) than in conventional TGA (14 g/m2). In situ mass measurements  show repeatability  during UHTC oxidation  testing  using  thermogravimetric  analysis where  consistent mass  gain  is  observed when similar specimen sizes are used. Parabolic rate constants determined using DNE-TGA from 1000°C-1300°C range from 0.037 to 0.0852 g2/m4\\x01s and from 1400°C-1600°C range from 0.257 to 2.22 g2/m4\\x01s in air over a 15 min hold. Model gas flow calculations showed that the range of 50-200 mL/min does not  influence the boundary  the gas velocity in  layer and does not significantly affect of ZrB2 + SiC in the TGA furnace. tional and DNE-TGA, we found that because specimens are  the oxidation behavior  In comparing  conven heated in air  continuously during  conventional  testing,  the  oxide layer is more protective (less B2O3 remains) because of reduced oxygen diffusion due to oxygen traveling through a  SiO2-rich  layer  that  results  in  less mass  gain  and  thinner  \\x0c', 'oxide  scales. The basic oxide  scale  structures are  similar  in  both methods where  specimens  exhibit  outer  glass  layers,  intermediate  layers,  and  inner  SiC-depleted  regions. How ever, the outer oxide scale 1600°C has ZrO2 particles more B2O3 remaining, while conventional  formed  during DNE-TGA at  across  its  length and likely has  the oxide  scale  formed during  does  not  exhibit ZrO2 layer or a separate SiC-depleted region. By using DNE-TGA to isothermally oxidize ZrB2 + SiC, we are able to determine that SiC in ZrB2 + SiC begins oxidation at ~1300°C.  particles  in  the  outer  Acknowledgments  The Air Force of Scientific Research (AFOSR) Young Investigator Program  Award, under grant number FA9550-10-1-0189,  supports DNE-TGA testing  capabilities. The National Science Foundation Early Faculty Career Award  under Division of Materials Research 0954110 supports material processing of  ceramics  using  spark  plasma  sintering. The Air Force Office  of  Scientific  Research Multidisciplinary University Research  Initiative  (AFOSR-MURI)  Program Award under grant number FA9550-1-1-0563 and Sub-award Num ber A001650204 supports gas flow and material  interaction modeling efforts.  The authors thank Fangyuan Gai of  the University of Arizona for performing  grain size measurements and Ronald E. Loehman for  technical  review of  the  manuscript.  Supporting Information  Additional  Supporting  Information may  be  found  in  the  online version of this article:  Fig.  S1.  Velocity  flow maps  from conventional  TGA  method using a gas flow of 50 mL/min.  Fig. S2. Conventional TGA mass measurements as a func tion of time showing dependence on surface area, where smaller specimens exhibit more mass gain at 1600°C. Fig. S3. EDS maps of (a) O, (b) Zr, and (c) Si of conven1600°C in  tional TGA oxidation  at  air  for  15 min  from  Fig. 3(c).  Fig. S4. 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},{
  "_id": 89,
  "PDF": "High-temperature oxidation and plasma torch testing of MoSi2–HfB2–MoB ceramics with single-level and two-level structure.pdf",
  "Text": "['Corrosion Science 158 (2019) 108074  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  High-temperature oxidation and plasma torch testing of MoSi2-HfB2-MoB ceramics with single-level and two-level structure  T  A.Yu. Potanin⁎, S. Vorotilo, Yu.S. Pogozhev, S.I. Rupasov, P.A. Loginov, N.V. Shvyndina, T.A. Sviridova, E.A. Levashov  National University of Science and Technology “MISiS”, Leninsky prospect 4, Moscow 119049, Russia  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ceramic material Combustion synthesis Oxidation Kinetics Mechanism Structure  Influence of the composition and structure of the heterophase ceramics in system MoSi2-HfB2-MoB on the thermal conductivity, kinetics and mechanisms of oxidation at 1200 °C and 1650 °C, including the plasma torch testing at 2000 °C, was investigated. Dense ceramics with heterogeneous single-level structure (SLS) and twolevel structure (TLS) were obtained by the hot pressing of powders produced by self-propagating high-temperature synthesis (SHS) in combustion mode. TLS ceramics have a lower thermal conductivity as compared to the SLS ceramics with similar elemental composition. Addition of hafnium diboride leads to the increase of mass change during the oxidation due to the formation of HfO2 and HfSiO4, which are denser than SiO2. In the case of the SLS ceramics, at 1200 °C a two-layered oxide film is formed. The upper layer is comprised of amorphous Hfdoped oxide layer, and the lower layer is crystalline crystoballite α-SiO2. TEM investigation of samples oxidized at 1650 °C revealed the formation of HfSiO4 precipitates in the α-SiO2 matrix. In the case of the TLS ceramics, regardless of oxidation temperature two-layered oxide film consists of upper SiO2 and lower HfSiO4. TLS ceramics demonstrated the highest oxidation resistance under the plasma torch and kept its structural integrity during 180 s at 2000 °C.  1.  Introduction  Transition metals silicide (TMS)-based composites are prospective as high-temperature oxidation and corrosion-resistant structural materials and coatings. They can withstand the mechanical loads at temperatures up to 1700-2000 °C in aggressive environments [1-3]. Among these materials, MoSi2-based composites attract special interest. In particular, they are applied as the heating elements in high-temperature industrial furnaces and exhaust systems of internal combustion engines [4-7]. Kanthal® (Sweden) in the current leader in the production of MoSi2based heating elements and produces elements capable of withstanding multiple heating up to 1850 °C in oxidative, neutral and reductive atmospheres [8]. In addition to the high oxidation resistance, advantages of MoSi2 include the high melting point (2030 °C), thermal shock resistance, and lower density (6.26 g/cm3) as compared to the Niand Cobased superalloys [9,10]. MoSi2-based ceramics are widely available, non-toxic and relatively cheap; moreover, they possess high electric and thermal conductivity, enabling the electro-discharge machining of complex ceramic parts.  The high-temperature oxidation resistance of MoSi2 is based on the formation of glassy SiO2-based protective film, which is nearly impermeable to gasses [11]. However, at the temperatures above 1800 °C the formation and evaporation mono-oxide SiO causes the degradation of the SiO2 film and drastically decreases the oxidation resistance of MoSi2. Moreover, an extensive industrial application of MoSi2 as the hightemperature structural material is hindered by a number of its inherent drawbacks. These include brittleness at T < 1000 °C, low creep resistance at T > 1200 °C, rapid oxidation in the air at 400-800 °C due to “pesting” phenomenon, and high discrepancy between the coefficients of thermal expansion (CTEs) of MoSi2 (8.3 × 10−6 К-1) and SiO2 (0.54 × 10−6 К-1), which leads to a decrease of the thermal cycling resistance below 900 °C [12-14]. Alloying by oxides Y2O3, La2O3, Al2O3, ZrO2, HfO2, borides TiB2, ZrB2, MoB, HfB2 and carbides SiC, ZrC, HfC allows one to enhance the performance of MoSi2-based materials [7,15-17] and optimize the protective properties of the oxide films [18]. CTEs of the transitional metal oxides are close to the CTE of MoSi2: for ZrO2 - 10.5 × 10−6 К-1, HfO2 - 6.0 × 10−6 К-1, TiO2 - 7.14 × 10−6 К-1. Lower CTE mismatch  ⁎ Corresponding author. E-mail address: a.potanin@inbox.ru (A.Y. Potanin).  https://doi.org/10.1016/j.corsci.2019.07.001 Received 11 April 2019; Received in revised form 4 July 2019; Accepted 6 July 2019  Available online 09 July 2019 0010-938X/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  mitigates the thermal stresses in the oxide film during the thermal cycling. However, such alloying often leads to the formation of eutectics and decreases the melting point and interval of working temperatures. In the MoSi2-MeIVB2 systems the eutectic reactions occur at 1880, 1900 and 1940 °C at TiB2, ZrB2 and HfB2 content equal to 15, 10, 7 at%, correspondingly [19]. For the oxide systems SiO2-TiO2, SiO2-ZrO2 and SiO2-HfO2 the eutectic melting temperatures are 1550, 1687 and 1695 °C [20]. The study of oxidation resistance of MoSi2 alloyed by 30 at% TiB2, ZrB2, HfB2 revealed that in the 500-1500 °C interval the composite MoSi2-HfB2 has the highest oxidation resistance. After 40 h of thermal cycling, mass growth at 1200 °C was 0.8 mg/cm2, at 1500 °C 1.9 mg/cm2 [16]. At T = 1000-1200 °C HfB2 undergoes rapid selective oxidation, and hafnia is the main oxidation product. At T = 1200-1400 °C the formation of SiO2-based film inhibits the selective oxidation of HfB2; correspondingly, the oxidation kinetics changes from the parabolic to the logarithmic-parabolic [21]. MoSi2-HfB2 composite demonstrated 6.5 times lower mass growth after 30 h of oxidation at 1400 °C as compared to MoSi2-ZrB2 [16]. Among the borides, MoB is another prospective alloying additive [15,22,23], which provides improved oxidation resistance and electric conductivity for MoSi2-MoB composites. A borosilicate glass film with high protective abilities is formed on the surface of the composites during oxidation [14,23]. Best oxidation resistance was achieved for the composition MoSi2-10 at% MoB [23]. To further augment the properties of MoSi2-based ceramics, optimization of the microstructure is pivotal. In this aspect, self-propagating high-temperature synthesis (SHS) is a particularly promising approach [24,25]. Variation of the SHS scheme allows one to produce the materials with similar phase compositions, but drastically different microstructures [26]. SHS was successfully employed to produce the MoSi2-HfB2-MoB ceramics with single-level structure (SLS) as well as two-level structure (TLS). In TLS composites, relatively large MoSi2 and MоВ grains (d = 10-15 μm) are surrounded by the layers of elongated HfB2 grains with the width up to 400 nm and length up to 2 μm [26]. TMS-based TLS composites are characterized by increased mechanical properties [26-28] and are promising for high-temperature applications [29]. However, the comparison between the oxidation mechanisms for SLS and TLS composites was not yet made. SHS allows the control of the grains size of individual phases in TLS composites by adjustment of the reaction mixture’s composition. For example, it is possible to increase the grains size of MoSi2 and decrease the grains size of HfB2. Thus, the microstructure of composite can be optimized to attain the best possible combination of mechanical properties and high-temperature oxidation resistance. The goal of this work was to investigate the influence of the comthe ceramics MoSi2-HfB2-MoB on the position and microstructure of thermal conductivity, mechanisms and kinetics of oxidation in furnace static conditions at 1200 °C and 1650 °C as well as in plasma torch dynamic conditions at 2000 °C.  2. Experimental procedure  In the investigated Mo-Hf-Si-B system, the composition of the combustion products was calculated according to the equation (100-X)∙ (90MoSi2+10MoB) + X∙HfB2, where X = 0 and 34 at% (40 wt%). Reaction mixtures were prepared using the following powders: Mo (purity 99.95%, particles size 2÷10 μm), Hf (purity 98.5%, particles size 30÷180 μm), Si size 2÷45 μm), B (purity 99.9%, particles size 1÷5 μm). The amorphous (purity 94.0%, particles calculated specimen’s elemental compositions are provided in Table 1. Dense ceramics were produced by hot pressing (HP) of the SHS powders using the DSP-515 SA installation («Dr. Fritsch», Germany) in vacuum at temperature 1600 °C, heating rate 10 °C/min, pressure 35 MPa and dwelling time 10 min. The raw SHS powders were produced according to two schemes  2  Table 1  Schemes of production of composite powders MoSi2-HfB2-MoB and calculated elemental compositions.  X, %  Scheme of production of composite SHS powder  Specimen  Composition, at%  X = 0 X = 34 X = 34  1 (individual synthesis) 1 (individual synthesis) 2 (co-synthesis)  X0 X34(1) X34(2)  Mo  35.0 23.2  Hf  -  11.3  Si  B  60.0 39.7  5.0 25.8  [26]:    In Scheme 1, products of the combustion of reaction mixtures Mo-Si-B (X = 0) and Hf + 2B were milled and mixed in a ball mill. 34 at% HfB2 were added (Table 1); In Scheme 2, products of the combustion of Mo-Hf-Si-B mixture (X = 34) were ball milled until the micron-sized MoSi2-HfB2-MoB composite powders were produced (Table 1).  Heat capacity (Cp) was measured by the differential scanning calorimetry using the DSC 404 C Pegasus installation (\"NETZSCH\", Germany). Platinum-rhodium crucibles with the alumina inserts and sapphire specimen with a known heat capacity (etalon samples) were used for measurements. Measurements were conducted under the argon flow of 50 ml/min. Thermal diffusivity (a) was measured by a laser flash method in the temperature diapason 20-400 °C under the argon flow (50 ml/min) on the NETZSCH LFA 457 MicroFlash installation (Germany). Thermal conductivity (λ) was calculated as λ = a·Cp·ρ. Density (ρ) was measured using the Archimedes method on analytic scales Sartorius ME235 (not less than 10 measurements in air and distilled water). Static oxidation experiments were conducted in air using the 10 × 10 × 5 mm3 specimens. Specimens were polished on «Rotopol21» (Struers, Denmark) and ultrasonicated in the isopropyl spirit. Oxidation at 1200 °C was conducted in the muffle furnace «SNOL 2.3 1.8/10» (АВ «UMEG-GROUP», Lithuania). Oxidation rate was estimated by the mass increase after 0.25; 0.5; 0.75; 1; 2; 3; 4; 5 h and further every 5 h until 30 h. The weight increase was measured with the accuracy of the 10−4 g. Static oxidation at 1650 °C was conducted in the «LHT 04/17 SW» furnace (Nabertherm, Germany) according to the following regimen: heating up to 1650 °C during 0.75 h; dwelling during 0.5 h; cooling with the furnace down to room temperature. The overall duration of the experiment was 5 h. During the experiments at 1200 °C, specimens were introduced in and extracted from the hot furnace, whereas during the experiment at 1650 °C specimens were introduced and extracted at the room temperature, while the heating and cooling were done with the furnace. Therefore, at the relatively low temperatures (400-800 °C) the oxidation occurred in the non-preferable “pesting” mode. Plasma torch dynamic oxidation tests were conducted on the UPIM200 installation under the following condition: plasma-forming gas flow 60 l/min (pressure 0.35 MPa), compressing gas flow 100 l/min (pressure 0.1 MPa), arc current 350 А, arc voltage 260 V, distance between the test sample and plasmatron nozzle 60 mm. Sample’s frontal temperature during the tests was measured using Termocont-TTs5C6M pyrometer. To measure the temperature field on the surface of samples and the change of the geometry of sample (linear erosion) we used the thermal imaging system «Tandem VS415». Using the optical and digital filters on the system, one can measure the linear erosion rate with the error margins ≤ 0.1 mm. the Gibbs energy (ΔGT) of Calculation of the possible oxidation reactions was performed on the online calculator FACT, developed at Ecole Polytechnique and McGill University, Canada (http://www.crct. polymtl.ca/FACT/). XRD analysis of  initial and oxidized samples was  the  conducted  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 1. Microstructure of the HP specimens with the compositions X0 (a), X34(1) (b, c) and X34(2) (d, e).  using monochromatic Cu Kα-radiation in the interval 2θ = 10÷110°. The microstructural was studied on a Hitachi S-3400 N scanning electron microscope (SEM) equipped with a NORAN energy dispersive Xray spectrometer (EDS). Quantitative information on the composition of the phases or the structural components were obtained by the EDS method, which was carried out at an accelerating voltage of 5÷20 kV. Individual samples were investigated by transmission electron microscope (TEM) JEM 2100 (JEOL, Japan). Electron-transparent specimens for TEM were prepared using the focused ion beam (FIB) method on the FEI Quanta 200 3D FIB instrument (FEI, Hillsboro, OR, USA).  3. Results  3.1. Microstructure, phase composition and thermal conductivity of consolidated ceramics  The microstructure of the HP specimens is provided in Fig. 1. In the X0 ceramic, MoSi2 (≤15 μm) and MoB (5-10 μm) grains are present. 34(1) ceramic (Fig. 1b, c) in addition to MoSi2 and MoB contains faceted HfB2 grains (4-10 μm) and a small number of HfSiO4 inclusions (≤4 μm). In both X0 and X34(1) ceramics, all phases have comparable size and are randomly distributed in the structure; hence these ceramics are denoted as single-level structures (SLS). 34(2) ceramic has a twolevel structure (TLS) (Fig. 1d, e) - relatively large MoSi2 grains (≤10 μm) are surrounded by the 2-4 μm wide HfB2 layers, which are comprised of elongated HfB2 grains with diameter 0.3 μm and length 0.5-1 μm (Fig. (≤3 μm) 1e). MoB grains are also present in the  3  \\x0c', 'A.Y. Potanin, et al.  Table 2  Phase composition of HP ceramics.  Phase (Structure type)  MoSi2 (tI6/2)  Corrosion Science 158 (2019) 108074  α-MoB (tI16/2)  HfB2 (hP3/4)  HfSiO4 (tI24/3)  β-MoB (oC8/2)  Specimen  Wt%  X0  X34(1)  X34(2)  90  48  50  Lattice parameter, nm a = 0.3204 c=0.7846 a = 0.3204 c=0.7847 a = 0.3204 c = 0.7843  Lattice parameter, nm a = 0.3117 c = 1.6936 a = 0.3117 c=1.6910  Wt%  10  9  -  Wt%  -  33  36  Lattice parameter, nm  a = 0.3137 c=0.3471 a = 0.3130 c=0.3458  Wt%  -  10  5  Lattice parameter, nm  a = 0.6577 c = 0.5972 a = 0.6571 c=0.5966  Wt%  Lattice parameter, nm  -  -  9  -  structure, usually in conjunction with HfB2 grains (Fig. 1e). This is related to the stages of crystallization in the combustion front of Mo-Hf-Si-B mixture: the most refractory HfB2 forms first, followed by crystallization of MoB on the surface of HfB2 [26]. Table 2 provides the phase composition of the samples. Main phases are the low-temperature tetragonal modifications of MoSi2 and α-MoB. The appearance of high-temperature rhombohedral β-MoB in the TLS X34(2) ceramic is related to reaction mixture higher combustion temperature (Tc = 2323 К) [26]. In both SLS X34(1) and TLS X34(2), HfB2 and HfSiO4 are present in similar proportion. Some physical and mechanical properties of ceramics are provided in Table 3. TLS specimen X34(2) demonstrated the highest hardness (HV) and lowest residual porosity (Pres). Interestingly, its thermal conductivity was the lowest among the investigated specimens (Table 3), despite the relatively high thermal conductivity of HfB2. To assess the influence of microstructure on the thermal conductivity of SLS and TLS MoSi2-HfB2-MoB composites, we used the relatively simple Brick Layer Model [30-32]. Calculated thermal conductivity for X34(1) and X34(2) was 36.8 and 28.4 W/m·K, which accounts for the difference in the HfB2 morphology in respective ceramics (Fig. 1) and coincides reasonably with the experimental measurements (Table 3).  3.2. Static oxidation of HP ceramics at 1200 °C  3.2.1. Oxidation kinetics The kinetic curves of the mass change during the static oxidation at 1200 °C is provided in Fig. 2. The shape of curves suggests that the oxidation is governed by semi-parabolic law: oxidation film growth rapidly at the beginning of oxidation, then the process is gradually inhibited. Within the first hour of oxidation (zone I) all samples demonstrate a linear mass increase. At this stage, the SLS specimen X0 the minimal mass gain (0.306 mg/cm2). After 1 h, demonstrates the curve for X0 becomes nearly horizontal, signifying the inhibition of oxidation. Specimens X34(1) and X34(2) demonstrate similar oxidation behavior. The discrepancy between their absolute values of mass gain increases during the first 10 h (zone II); the effect of further oxidation is negligible (zone III). Addition of HfB2 to the MoSi2-MoB ceramic increases the weight gain during the oxidation due to the formation of hafnium oxide and silicate, which have higher specific weights as compared to silica (SiO2 - 2.36 g/cm3, HfO2 - 9.68 g/cm3, HfSiO4 -  Table 3  Physical and mechanical properties of HP ceramics.  Specimen  ρ, g/cm3  Pres, % [26]  HV10, GPa [26]  a,* mm/s2  X0 X34(1) X34(2)  5.90 7.31 7.51  3.5 3.3 0.8  9.2 13.4 17.6  9.66 10.43 8.31  Cp,* J/ g·K  0.50 0.44 0.42  λ,* W/ m·K  28.50 33.55 26.21  *  data for 400 °C.  6.97 g/cm3) [21]. Although the TLS X34(2) ceramic has higher relative density than SLS X34(1) and nearly identical phase composition, its weight gain is somewhat higher than for X34(1). The discrepancy in weight gains is probably caused by the more intense oxidation of the MoB and HfB2 grains in TLS X34(2) ceramic due to their smaller size and particular arrangement (Fig. 1b, c).  3.2.2. The microstructure of oxide films Fig. 3 demonstrates the fractures of the near-surface zone of specimens oxidized at 1200 °C during 0.5 and 30 h. For all compositions, oxidation during 30 h leads to the formation of two-layered oxide films with vastly varying structure and phase composition. A continuous borosilicate film is formed on the surface of all the specimens. The the oxide films was 2-4 μm after 0.5 h; after 30 h, width of the oxide films were 5-15 μm wide. Local inclusions of MoO2 are present in all films in accordance with [11,23]. For the SLS X0 specimen, 1-3 μm grains of Mo5Si3 can be identified on the interface between the ceramic and borosilicate oxide film after 30 min oxidation. Oxidation during 30 h leads to the formation of a continuous 2 μm thick Mo5Si3 layer (Fig. 3a-c). This interlayer forms due to the Si depletion in the near-surface zone of ceramic as a result of the SiO2 formation. The formation of Mo5Si3 is thermodynamically more preferable that the formation of MoO3, MoO2, and Mo (see Section 3.2.3). The microstructure of the oxide layers formed after 0.5 h on the surface of specimens X34(1) and X34(2) differs considerably. HfSiO4 content in X34(1) is several times lower due to relatively slower oxidation of coarser HfB2 grains; the oxide film is comprised of SiO2 grains surrounded by B2O3 interlayers. After 30 h, self-organized two-layered oxide films were formed on the surface of both X34(1) and X34(2) (Fig. 3f, j). On the surface of SLS X34(1) specimen, two oxide layers with different Hf content were formed. These layers are discernible on the SEM images due to high compositional contrast. The top layer was amorphous, 9-12 μm wide and characterized by higher Hf content than the 3-5 μm wide crystalline bottom layer (Figs. 3f, g, 4 ). In case of TLS X34(2) sample, 10 μm wide upper layer is formed of SiO2, whereas the lower 5 μm wide layer is comprised of HfSiO4 grains with the size of 1-3 μm. (Fig. 3j, k). Both layers are crystalline. HfSiO4 grains with 0.5-2 μm size embedded in the SiO2-B2O3 matrix are present on the surface of TLS specimen X34(2) after 0.5 h oxidation (Fig. 3i). Extensive formation of HfSiO4 during the oxidation is peculiar to the TLS X34(2) sample and is caused by the selective oxidation of fine HfB2 grains due to the effective grain boundary diffusion of oxygen. B2O3 interacts with SiO2 and forms (in the 417-456 °C interval) the eutectic melt SiO2-B2O3, which has high evaporation temperature and effectively blocks the access of oxygen towards the unoxidized material. Afterward, interdiffusion of SiO2 and HfO2 leads to the formation of HfSiO4. Due to the high wetting angle, B2O3 melt contracts the HfSiO4 grains together, leading to the densification of the oxide film [33]. However, at the temperatures above 1100 °C the B2O3 melt begins to evaporate [21]. The high specific density and poor  4  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 2. Kinetic curves of oxidation at 1200 °C.  wettability of HfSiO4 grains by SiO2-B2O3 melt lead to the formation of a dense, continuous HfSiO4 layer at the interface between the ceramic and SiO2 film. With the increase of the oxidation duration, the HfSiO4 grains are relocated to the surface of unoxidized ceramic forming the continuous HfSiO4 layer, and the upper SiO2 layer growth incrementally, hindering the further diffusion of oxygen into the ceramic. Correspondingly, the oxidation curve becomes nearly horizontal (Fig. 2). Fig. 4 provides the TEM images of the SLS X34(1) specimen after the oxidation. Two SiO2-based layers form on the surface of the specimen:  to the  layer, adjacent  amorphous upper layer and a crystalline lower unoxidized ceramic (Fig. 4a, b). According to EDS, upper layer contains 2.4 ± 0.7 at% Hf in addition to Si and O. Presumably, the dissolved Hf leads to amorphization of SiO2. Hf and Si are isovalent elements; therefore, Hf can be incorporated in the SiO2 lattice with the formation of a Six-1HfxO2 solid solution. Due to Hf’s higher atomic radius (0.167 nm) as compared to Si (0.132 nm), the Si-O-Si bond is stretched. A similar pattern was observed in [34,35], where the alloying of TiO2 by Zr led to the formation of Tix-1ZrxO2 solid solution and inhibited the transition of TiO2 from  Fig. 3. SEM images of the fractures of specimens X0 (a-c), X34(1) (d-g) and X34(2) (h-k) oxidized at 1200 °C during 0.5 h (a, d, e, h,  i) and 30 h (b, c,  f, g,  j, k).  5  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 4. TEM image of the near-surface zone of oxidized SLS specimen X34(1) (a); diffraction patterns and EDS of the oxide layers (b); dislocation network in the MoSi2 grain near the boundary with the oxide layer (c); HRTEM image of the phase boundary between MoSi2 and α-SiO2 and diffraction pattern from MoSi2 grain (d).  amorphous to the crystalline state (anatase and rutile) upon heating. The dissolution of Hf in the upper oxide layer of the X34(1) ceramic (Fig. 3f, g) instead of formation of HfO2 or HfSiO4 is related to the size of HfB2 grains in the ceramics. Due to their larger size (up to 15 μm), HfB2 grains in the SLS X34(1) ceramic oxidize slower than HfB2 grains (0.5-3 μm) in the TLS X34(2). Since the Hf possesses higher chemical affinity towards oxygen as compared to Si, self-organization of the twolayered oxide film occurs, with an upper layer containing the dissolved Hf. A large discrepancy between the surface energies of crystalline and amorphous silica [36] results in the formation of the high-energy phase boundary between the layers. This high-energy boundary prevents the into the bottom crystalline α-SiO2 layer and stabilizes diffusion of Hf the two-layered structure of oxide film. Fig. 4c demonstrated the MoSi2 grain with a high concentration of dislocations. Dislocations are presumably formed due to thermal stresses caused by temperature gradients and the discrepancy between the coefficients of thermal expansion of MoSi2 and SiO2. According to the high-resolution electron microscopy  (HRTEM)  (Fig. 4d), the phase boundary between MoSi2 and α-SiO2 is incoherent. Diffraction pattern (Fig. 4d) of MoSi2 grain, recorded along the zone axis [0-10], shows the body-centered tetragonal symmetry of MoSi2 according to the XRD data (Fig. 5).  3.2.3. Oxidation mechanism The formation of protective layer during the oxidation of MoSi2 can occur according to reactions (1-4) [23]:  5MoSi2(s) + 7O2(g) → Mo5Si3(s) + 7SiO2(s)(ΔG1200 °C = −4228 kJ/mol) (1)  2MoSi2(s) + 7O2(g) → 2MoO3(g) + 4SiO2(s)(ΔG1200 °C = −2870 kJ/mol) (2)  MoSi2(s) + 3O2(g) → MoO2(s) + 2SiO2(s)(ΔG1200 °C = −1494 kJ/mol) (3)  MoSi2(s) + 2O2(g) → Mo(s) + 2SiO2(s)(ΔG1200 °C = −1169 kJ/mol)  (4)  6  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 5. XRD patterns of the specimens oxidized at 1200 °C during 30 h.  Boron oxide B2O3(l) forms due oxidation of MoB and HfB2 according to reactions (5-7) [21,23]. B2O3 has a low melting temperature (450 °C) and forms a viscous melt on the surface of ceramics, hindering the oxygen diffusion into the material.  2MoB(s) + 3,5O2(g) → B2O3(l) + 2MoO2(s)(ΔG1200 °C = −1361 kJ/mol) (5)  2MoB(s) + 4,5O2(g) → B2O3(l) + 2MoO3(g)(ΔG1200 °C = −1242 kJ/mol) (6)  HfB2(s) + 2,5O2(g) → B2O3(l) + HfO2(s)(ΔG1200 °C = −1478 kJ/mol) (7)  Nano-precipitates of MoB can form under the silica layer as a result of reaction (8) between the MoSi2 and B2O3 [37,38]:  2MoSi2(s) + B2O3(l) + 5/2O2(g) → 2MoB(s) + 4SiO2(s)(ΔG1200 −1492 kJ/mol)  °C =  (8)  At higher temperatures, a layer of hafnium silicate HfSiO4 can be formed via the reaction (9) [21,39]:  HfO2(s) + SiO2(s) → HfSiO4(s)(ΔG1200 °C = −2363 kJ/mol)  (9)  HfSiO4 is similar to ZrSiO4 and likewise promotes the self-healing of pores and cracks in the SiO2-B2O3 oxide layer and acts as a diffusion barrier [40]. Moreover, HfSiO4 increases the crystallization temperature of the amorphous films, therefore increasing the oxidation resistance of the silicon-containing materials at high temperatures [41]. The melting point of HfSiO4 is 1750 °C. Presumably, in the MoSi2-HfB2-MoB system the borides begin to oxidize before MoSi2. As-formed B2O3 actively evaporates, promoting  the interaction of borides with oxygen. After the oxidation of MoSi2 begins, the process is impeded by the formation of borosilicate glass (10):  xSiO2(s) + yB2O3(l) → xSiO2∙yB2O3(s,l)  (10)  Since the HfB2 has a small composition variability (65.5÷67.7 at% B [42]), the following reaction can occur after the oxidation of MoSi2 with the formation of SiO2 (11):  HfB2(s) + SiO2(s) → Hf1-xB2 + Six-1HfxO2(s)  (11)  XRD analysis of the surface of samples oxidized during 30 h at 1200 °C is provided in Fig. 5. Besides the initial phases MoSi2, MoB, the oxide phases α-SiO2, MoO2 HfB2, HfSiO4, and HfO2, formed in correspondence with reactions (1-9), were identified. α-SiO2 is a hightemperature modification of quartz (cristobalite), stable at high temperatures and low pressures [43]. HfO2 has a low-temperature modification, stable up to 1650 °C [44]. In the SLS specimen X0, Mo5Si3 was formed corresponding to the reaction (1). During the oxidation, HfSiO4 content increases from 5 to 10% (in the initial ceramic) up to 15-30% due to reactions 7 and 9. Also, in the TLS specimen X34(2) high-temperature rhombohedral modification transforms into low-temperature tetragonal α-MoB. The oxide films formed on the surface of X34(1) and X34(2) show a noticeable difference in the orientation of HfSiO4 grains. In case of X34(1) ceramic, the intensity of (101), (200) and (112) peaks located at 2θ = 20.1, 27.1, 35.7° corresponds to the polycrystalline HfSiO4. In case of X34(1) ceramic, all HfSiO4 forms within the ceramic itself during the sintering; no HfSiO4 was present in the oxide film (Table 2, Figs. 3, 4). For X34(2) ceramic, some of the HfSiO4 grains are located within the ceramic itself (Table 2), and others in the HfSiO4 layer of the  7  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 6. SEM images of the near-surface zone of  fractured specimens X0 (a), X34(1) (b, c) and X34(2) (d, e) oxidized at 1650 °C during 30 min.  two-layered oxide film (Fig. 3). For oxidized X34(2), the (200) peak of (2θ = 27.1°) HfSiO4 is approx. 2 times more intense than (101) and (112) peaks. Since the XRD patterns of HfSiO4 within X34(1) and X34(2) before oxidation were similar, the pronounced increase in the intensity of (200) peak after oxidation suggests the preferential (200) orientation of the HfSiO4 layer in the oxide film.  3.3. Static oxidation of HP ceramics at 1650 °C  3.3.1. The microstructure of the oxide films Fig. 6 provides the microstructure of fractures of surface regions of specimens oxidized at T = 1650 °C during 0.5 h. The thickness of the SiO2 layer for SLS specimen X0 is 4-5 μm (Fig. 6a). Some SiO2 grains 1-2 μm in size are located under the oxide layer, mainly on the boundary between MoSi2 grains and in microcracks. Defects, such as pores and microcracks, and grains boundaries are most susceptible to oxidation [45].  On the surface of TLS X34(2) ceramics at 1650 °C as well as at 1200 °C a self-organized two-layered oxide film is formed: upper 5-6 μm thick continuous layer consists of SiO2, while lower 6-8 μm thick layer comprised of HfSiO4 grains 2-4 μm in size (Fig. 6d, e). On the surface of SLS X34(1) specimen, a one-layered SiO2-based 8-10 μm thick film (Fig. 6b), containing fine HfSiO4 precipitates (Fig. 6c). The following mechanism for the formation of HfSiO4 has been formulated. The presence of oxides (including B2O3) facilitates the dissolution of HfO2 in SiO2 [46,47]; similarly, the presence of B2O3 during the initial stage of oxidation supposedly enhances the solubility of HfO2 in the HfO2-SiO2-B2O3 melt. When the specimen reaches the isothermal oxidation temperature (1650 °C), B2O3 and SiO2 sublimate as SiO and B2O2, resulting in the oversaturation of the residual oxide melt by Hf. As a result, nano-scale HfSiO4 precipitates are formed. Previously, ZrO2 nano-precipitates were observed in the SiO2 oxide film after the high-temperature oxidation of ZrB2-MoSi2 [38] and ZrB2-SiC [48] ceramics.  8  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 7. TEM image of the near-surface zone of SLS X34(1) specimen (a) containing HfSiO4 precipitates (b); diffraction pattern (c), EDS analysis (d) and HRTEM image (e) of the precipitate.  Table 4  Phase composition of the surface of specimens after the oxidation at 1650 °C during 30 min.  Specimens  X0 X34(1) X34(2)  Identified phases (structure type)  α-SiO2 (tP12/1), MoSi2 (tI6/2), α-MoB (tI16/2) α-SiO2 (tP12/1), MoSi2 (tI6/2), α-MoB (tI16/2), HfB2 (hP3/4), HfSiO4 (tI24/3), HfO2 (mP12/3) β-SiO2 (cF24/3), MoSi2 (tI6/2), α-MoB (tI16/2), HfB2 (hP3/4), HfSiO4 (tI24/3), HfO2 (mP12/3)  Weight gain, mg/cm2  0.329 1.226 1.440  Fig. 9. Test specimen in the holder.  (Fig. 7c) and EDS analysis (Fig. 7d), these precipitates were identified as HfSiO4 with body-centered tetragonal lattice. The lattice parameters, measured by diffraction pattern (a = 6.957 Å, c = 6.094 Å), demonstrated a slight deviation from the standard values (a = 6.580 Å, c = 5.980 Å, ICDD card № 77-1759). This deviation might have been caused by the dissolution of some hard-to-identify elements (such as B), measurement error, non-homogeneous composition or defects in the crystal lattice of HfSiO4. HRTEM (Fig. 7e) demonstrated that the submicron precipitates have no subgrains and are defect-free.  Fig. 8. XRD patterns (low-angle area) of the samples oxidized at 1650 °C during 30 min.  The oxide film formed on the SLS ceramic X34(1) was investigated by TEM (Fig. 7). Elongated dark-grey precipitates are distributed homogeneously in the SiO2 matrix and have the length of 300-400 nm and width of 100-200 nm (Fig. 7b). By the electron diffraction pattern  9  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Fig. 10. Temperature field distribution at 150 s for samples X0 (a), X34(1) (c), and X34(2) (e); frontal temperatures during the test for specimens X0 (b), X34(1) (d), and X34(2) (f).  3.3.2. Oxidation mechanism and kinetics Although the total duration of oxidation at 1650 °C was 6 times lower than at 1200 °C, total weight gain was 20-30% higher. Similar to oxidation at 1200 °C, the TLS specimen X34(2) had the highest weight gain (1.44 mg/cm2) after oxidation at 1650 °C (Table 4). The results of XRD analysis of the surface of specimens oxidized during 30 min at 1650 °C are presented in Table 4. For the SLS X0 specimen, the Mo5Si3 is absent, despite being present  after the oxidation at 1200 °C. According to [49,50], this phase does not form during the oxidation at T > 1400 °C. Presumably, after the formation of a continuous SiO2 layer on the surface of the specimen, the formation of Mo5Si3 according to reaction 1 is arrested. At higher temperatures, Mo5Si3 begins to oxidize by reaction 12:  2Mo5Si3(s) + 21O2(g) → 10MoO3(g) + 6SiO2(s)(ΔG1650 °C = −5166 kJ/ mol) (12)  10  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  Samples X0 and X34(1) experienced a similar heating rate. During the first 95 s the temperature increased up to 1670-1710 °C. In the next 5-7 s, temperature has risen drastically to 1780-1800 °C and then remained nearly constant until the 150 s mark. The intermediate temperature rise was presumably related to the melting and removal of the oxide layer, which was formed during the first period (0-95 s). The newly exposed materials surfaces had higher catalytic activity, which was the cause of the rapid increase of the surface temperature. The rise of surface temperature after the 150 s mark is presumably caused by the MoSi2 = 2030 °C) and resulting intense erosion of melting of MoSi2 (T m material. In the case of SLS ceramics X0 and X34(1), the tests were canceled after 160 s due to the destruction of the samples. However, the TLS X34(2) ceramic demonstrated noticeably better oxidation and erosion resistance. Sample experienced relatively stable heating up to 2000 °C during the first 75 s with minimal erosion and surface melting. During the 75-180 s period the sample was exposed to the harshest testing conditions but retained its form and structural integrity. This enhanced oxidation and erosion resistance can be attributed to the peculiar structure and composition of the formed oxide layer. XRD patterns of the samples X34(1) and X34(2) after plasma torch tests are provided in Fig. 11. SLS sample X34(1) consists of the MoSi2, MoB, HfB2, and HfO2, i.e., the phases which were present in the specimen before the testing. Phase composition of X34(1) corroborates the assumption that the oxides were completely removed from the specimen’s surface by the plasma flow. However, XRD pattern for TLS sample X34(2) shows the presence of intense SiO2 and HfSiO4 peaks, suggesting that the protective oxide film was retained during the test. After the dynamic plasma torch oxidation, HfSiO4 peaks show no signs of preferential orientation as opposed to the oxidation at 1200 °C, presumably due to the intense recrystallization during the plasma torch testing.  3.4.2. The microstructure of the oxide layers Fig. 12 demonstrates the microstructure (polished transverse cut) of TLS X34(2) specimen after the plasma torch test. Continuous SiO2-HfSiO4 two-layered film was formed on the surface of the specimen and protected it from oxidation and abrasion during the test. The self-organizing film consists of an upper 20-25 μm wide SiO2 layer and a lower 15-20 μm wide HfSiO4 layer. The mechanism of formation of HfSiO4 layer was described in Section 3.2.2. Upper SiO2 oxide layer contains numerous spherical HfSiO4 precipitates with the size of 30-100 nm. The mechanism of formation of HfSiO4 precipitates is similar to the one described in Section 3.3.1; the difference in the form of precipitates might be attributed to spheroidization of precipitates due to the higher testing temperature. Underneath the HfSiO4 layer, 1 μm MoB grains were formed according to the reaction (8). Interestingly, despite having the highest weight gain during the static oxidation at 1200 and 1650 °C, TLS X34(2) ceramic showed far better resistance towards the dynamic oxidation under the plasma  Fig. 11. XRD patterns of the ceramics X34(1) and X34(2) after the plasma torch tests.  During the oxidation of TLS sample X34(2), high-temperature rhombohedral β-MoB transforms tetragonal αinto low-temperature MoB. Tetragonal α-SiO2 was identified on the surface of SLS samples X0 and X34(1), whereas cubic β-SiO2 was present on the surface of TLS X34(2). β-SiO2 forms at 1470 °C and is stable up to the melting point (1728 °C). Upon cooling to 270 °C the reverse transformation of βSiO2 into α-modification occurs, leading to the 5% increase of oxide’s density and formation of cracks in the oxide layer [51]. However, in the case of TLS specimen X34(2) the β-SiO2 was presumably stabilized by HfO2 and HfSiO4. The inhibition of β→α transformation by some oxides (Na2O, CaO, Al2O3, ZrO2) was previously reported [52-54]. Fig. 8 demonstrates the XRD pattern of the oxidized samples in the angle interval 19÷23°, where the most intense SiO2 peaks are located. For the X34(2) specimen, the diffraction peak shifts towards lower angles. For all the Hf-containing specimens, (101) peak of is presented on the XRD patterns. This peak is the second most intense peak (after (200)) of the hafnium silicate HfSiO4.  3.4. Dynamic oxidation during plasma torch testing  3.4.1. Oxidation kinetics and temperature profiles of specimens Cylindric test samples (Fig. 9) with Ø = 20 mm and h = 12 mm were fastened in a carbon-carbon composite holder. Fig. 10 provides the temperature field distribution in specimens (Fig. 10a, c, e) and specimen’s frontal temperatures (Fig. 10b, d, f) during the plasma torch testing (180 s). For the SLS samples X0 and X34(1), cracking, partial melting and intense erosion were observed.  Fig. 12. Microstructure of the oxide film of TLS X34(2) ceramic after the plasma torch test with the magnification × 2 000 (a) and × 20 000 (b).  11  \\x0c', 'A.Y. Potanin, et al.  Corrosion Science 158 (2019) 108074  torch. This enhanced performance can be attributed to the peculiar twolevel microstructure of X34(2) ceramic, which affects the performance in two ways. First, the MoSi2 grains are encapsulated by HfB2 layers and are less affected by the direct impact of plasma flow. Second, the preferential oxidation of fine HfB2 grains leads to the formation of a twolayered oxide film with HfSiO4 nano-precipitates in SiO2 layer and continuous 10-20 μm thick HfSiO4 layer. The resulting two-layered oxide film is hard, strong, stable and refractory enough to withstand the thermal, mechanical and erosive impact of plasma flow. Hence, the TLS X34(2) ceramic, despite having the highest weight gain during stationary oxidation at 1200 and 1650 °C, vastly outperformed both SLS X0 and SLS X34(1) ceramics in the dynamic oxidation under the plasma torch. Understanding and control of the phase and structure formation of oxide films in single-level structured (SLS) and two-level structured (TLS) ceramics open new avenues for the optimization of the composition and structure of ceramics in regards to the intended working conditions. Application of SHS is particularly promising in this regard since it allows one to produce a wide spectrum of ceramics with similar phase composition but drastically different microstructures and properties [26-28].  4. Conclusions  1 The mechanisms of formation of various self-organized two-layered oxide films and HfSiO4 nano-precipitates are described and discussed with an emphasis on the role of single-level structure (SLS) or two-level structure (TLS) of MoSi2-HfB2-MoB ceramics. 2 At 1200 °C, oxidation of SLS and TLS ceramics produced two-layered oxide films. In the case of SLS ceramic, SiO2-based oxide film was top 9-12 μm wide Hf-doped amorphous comprised of silica layer and bottom 3-5 μm crystalline α-SiO2 layer. In the case of TLS crystalline α-SiO2, whereas ceramic, top layer consisted of the bottom layer was formed by HfSiO4 grains with preferential (200) orientation. 3 Precipitation of nanoscale HfSiO4 grains in the SiO2 matrix occurs at 1650 °C during static oxidation for the SLS ceramics and at 2000 °C during dynamic oxidation for the TLS ceramics. 4 During the static oxidation at 1200 and 1650 °C, the TLS ceramic demonstrated the highest weight gain due to preferential oxidation of fine HfB2 grains and formation of 5-8 μm wide crystalline HfSiO4 sublayer. However, in the case of dynamic oxidation under the plasma torch, this sublayer provided sufficient protection against plasma flow and defined the best performance of TLS ceramic under the plasma torch test.  Acknowledgment  The work was supported (Agreement No. 19-19-00117).  by  the  Russian  Science  Foundation  References  [3]  [1] H. Qiang, M. Chaoli, Z. Xinqing, X. 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},{
  "_id": 90,
  "PDF": "High-Temperature Oxidation at 1900 degrees C of ZrB2-xSiC Ultrahigh-Temperature Ceramic Composites.pdf",
  "Text": "['High-Temperature Oxidation at 19001C of ZrB2-xSiC Ultrahigh Temperature Ceramic Composites  Wen-Bo Han, Ping Hu,  w  Xing-Hong Zhang, Jie-Cai Han, and Song-He Meng  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, China  Oxidation of ZrB2-based ultrahigh-temperature ceramic composites containing 10, 20, and 30 vol% SiC was performed at 19001C for 1 h in air. ZrB2-20 vol% SiC exhibited the highest oxidation resistance at this temperature and formed a dense co herent oxide scale after oxidation, whereas a strong degradation  was observed for both ZrB2-10 vol% SiC and ZrB2-30 vol% SiC. In addition, cracks and spallation in the oxide scale were  also detected for the latter materials. The oxidation behaviors of ZrB2-SiC composites were investigated. The effect of SiC content was analyzed and oxidation models were proposed to de scribe the observed microstructures.  I.  Introduction  THE ZrB2and HfB2-based ultrahigh-temperature (UHTCs) are potential candidates for thermal protection materials in both re-entry and hypersonic vehicles because of  ceramics  their excellent and unique combination of high melting points,  good thermal shock resistance, and good ablation/oxidation resistance.1-12 These properties make UHTCs attractive for the  design of future hypersonic aerospace vehicles with features like  sharp leading edges and sharp nose cones. Such design features  could produce more agile vehicles that would open up a greater  range of hypersonic flight paths and re-entry trajectories. Re entry and hypersonic vehicles, regardless of their specific designs,  require control surfaces with sharp leading edges if they are to be  maneuverable at hypersonic velocities. Low-radius leading edges  are  subject  to much greater  aerothermal heating  than blunt  edges, such as those on the space shuttle orbiter, and these edges will reach temperatures that may exceed 20001C during reentry.13,14 The currently available thermal protection materials will  not survive under such extreme temperatures and new materials  are required for use in advanced thermal protection systems.  Oxidation resistance is a major issue in the development of  UHTCs for aero-propulsion and hypersonic flight applications.  The introduction of  second phases  (i.e., SiC, MoSi2) has ceeded in improving both the oxidation/ablation resistance and the mechanical properties of UHTCs.3-6,8,11,15 In particular, the  suc addition of Ta compounds significantly improves the oxidation resistance of ZrB2-SiC below 18001C, whereas it is detrimental to the performance of ZrB2-SiC at a higher temperature.16,17 In accordance to the previously reported results, a diboride matrix  composite that includes only SiC as a second phase is one of the  most promising compositions. However,  the oxidation mecha nisms of these materials are still not well understood, especially 19001C.  for  temperatures  above  In  addition,  the  oxidation  behavior of  the ZrB2-SiC composites from that at lower temperature. The purpose of this paper is to  is  significantly different  investigate the oxidation behaviors of ZrB2-SiC UHTC composites at 19001C in air. The effect of SiC content on the oxi dation resistance of  the materials was  investigated,  and the  proposed model was also discussed.  II.  Experimental Procedure  The samples used here for oxidation testing were fabricated from  commercial ZrB2 Research, Xi’an, China) and SiC (Weifang Kaihua Micro-pow (Northwest  Institute  for Non-ferrous Metal  der Co. Ltd., Shandong, China) powders. The powder mixtures  of ZrB2110 vol% SiC (ZS1), ZrB2120 vol% SiC (ZS2), and ZrB2130 vol% SiC (ZS3) were ball milled in ethanol for 8 h and dried in a rotating evaporator. Milled powders were then uniaxially hot pressed in a boron nitride-coated graphite die at 20001C  for 60 min under vacuum and 30 MPa of applied pressure.  Bulk density and theoretical density were evaluated using the  Archimedes method and the rule of mixtures, respectively. Sample coupons in the size of 2.0 cm \\x02 1.0 cm \\x02 0.35 cm were cut from the hot-pressed materials, and all surfaces were diamond polished to a 1-mm finish. Coupons were ultrasonically cleaned successively  in detergent, deionized water, acetone, and alcohol before expo sure. All samples were loaded into a slotted ZrO2 refractory brick and then exposed to 1-h oxidation in stagnant air at 19001C using  a bottom-loading furnace with zirconia heating elements, and the  results are repeatable. The experiments were carried out with a  Pt-Rh thermocouple and a two-color Raytek pyrometer (RAY MR1SCSF, Raytek Corp. Santa Cruz, CA), which covers a temperature range of 10001-30001C. MoSi2 elements were used to heat before the operation of the zirconia heating elements, and the  temperature was measured by the Pt-Rh thermocouple. The two color Raytek pyrometer was used to measure the temperature of  the samples.  X-ray diffraction (Rigaku, Dmax-rb, Tokyo, Japan) was used  to identify oxide phases present after exposure. Scanning elec tron microscopy (SEM; FEI Sirion, Eindhoven, Holland) along  with energy-dispersive spectroscopy (EDAX, Mahwah, NJ) was  used to characterize the composition and microstructure of the  surface and cross  section of  the samples after oxidation. The  different oxide layers were also investigated after removing the  surface layers by polishing parallel to the original surface. The  material  removal was monitored using optical microscopy so  that the desired region was reached.  III.  Results and Discussion  (1)  Density and Microstructure  Dense ZrB2-SiC composites were obtained. The bulk densities of the sintered ZS1, ZS2, and ZS3 billets were 5.80, 5.41, and 5.22 g/cm3, which correspond to relative densities of 100%,  98%, and 100%, respectively. Figure 1 shows SEM micrographs  of  the polished surfaces of  the ZrB2-SiC UHTC composites. The darker phase is SiC, and it appears to be uniformly dis persed in the lighter ZrB2 matrix. The microstructures of  the  N. S. Jacobson—contributing editor  This work was  supported by  the National Natural Science Foundation of China  (50602010),  the  Research  Fund  for  the Doctoral  Program of Higher  Education  (20060213031) and the Program for New Century Excellent Talents in University.  w  Author to whom correspondence should be addressed. e-mail: huping@hit.edu.cn  Manuscript No. 24445. Received March 21, 2008; approved July 22, 2008.  Journal  J. Am. Ceram. Soc., 91 [10] 3328 - 3334 (2008)  DOI: 10.1111/j.1551-2916.2008.02660.x  r 2008 The American Ceramic Society  3328  \\x0c', 'October 2008  High-Temperature Oxidation at 19001C of ZrB2-xSiC  3329  (a)  (b)  (c)  10 µm  10 µm  10 µm  Fig. 1.  Scanning electron micrographs of the surfaces of the ZrB2-SiC ultrahigh-temperature ceramics: (a) ZS1, (b) ZS2, and (c) ZS3.  composites are regular, and little porosity was observed in the  (2)  Microstructure Changes and Oxidation Properties  polished surfaces. Based on the high relative density and the lack  of any open porosity, porosity should not have a significant  effect on oxidation behavior.  Figure 2 shows the micrographs of ZrB2-based UHTCs containing 10, 20, and 30 vol% SiC after oxidation at 19001C for  1 h. The bright phase is ZrO2 and the dark phase is SiO2, as  (a)  (b)  100 µm  100 µm  (c)  Fig. 2.  Scanning electron micrographs of the surfaces of the ZrB2-SiC ultrahigh-temperature ceramics after oxidation at 19001C in air for 1 h: (a) ZS1, (b) ZS2, and (c) ZS3.  100 µm  \\x0c', 'shown in Fig. 2. A continuous SiO2 layer was found to form on the surface of the oxide scale for ZS2, while a discontinuous  SiO2 layer was formed at the outside scale for ZS1 and ZS3. A number of holes were also observed on the surface of ZS1 and  ZS3. The formation of the holes was attributed to the formation  of  large amounts of high-pressure gaseous oxidation products  (i.e., B2O3, CO, and SiO). Moreover, the oxide scale was flaky and brittle to the touch. No boron was detected in the surface  layer of all the samples in the present situation. Cross-sectional SEM micrographs of ZS3 oxidized at 19001C  in air for 1 h are shown in Fig. 3. There are four distinct layers in  the oxide scale: (1) a silica-rich outer layer, (2) a subscale of crys talline zirconia, containing little silicate,  (3) a partially depleted  zirconium diboride layer along with a completely depleted layer in  SiC grains, and (4) an unaltered material. The thickness of  the  first and second layer is far thinner than the third one. Appar ently, the present condition is favorable for the active oxidation of  SiC, which accounts for the generation of this layer with a rela tively high thickness. Further analysis shows that the initial con tinuous ZrB2 matrix was (Fig. 3b), which has not been reported in the previous ture.3,6,8,18,19 Examination by EDS (not  transformed  into discrete  structure  litera shown)  indicated the  presence of zirconium and boron as  the primary elements  that  corresponded to ZrB2. Figure 4 shows SEM micrographs of the subside scale of ZS3  after successfully removing the outside scale. This is an interface  between layer 2 and layer 3, which contains  two distinctive  structures, namely, an intact oxide scale and a separate oxide  scale as shown in Fig. 4(a). A discrete ZrO2/ZrB2 boundary was observed in Figs. 4(a) and (c). As can be seen, the ZrB2 particles were partially consumed and ZrB2 structure, initially connected, did not exist (Fig. 4(b)), which is consistent with the morphology  of the oxide scale in Fig. 3(b). The oxide products did not adhere  to the unaltered material in this region (Fig. 4(b)). However, the  oxide products in the other region were adhered to the unreacted  material  (Fig. 4). The dark silica-containing phases were em bedded  in  bright  phases,  which  were  identified  as  ZrO2 (Fig. 4(d)). The ZrO2 was sintered into an integrated structure at this temperature.  A number of pores were detected in the oxide scale of ZS1 after oxidation at 19001C, as shown in Fig. 5(a). However, no  SiC-depleted zone was detected in this material at this temper ature (Fig. 5(b)). The formation of the SiC-depleted layer in the  ZrB2-SiC system depends not only on the surrounding conditions of pressure and temperature but also on the structure dis tribution of SiC in the ZrB2 matrix. SiC particles were dispersed in the ZrB2 matrix and the SiC content was too low to form interconnectivity in ZS1. Therefore, no SiC-depleted layer was  formed because the diffusion of oxygen in ZrB2 was very low. Figure 6 shows the cross-sectional micrographs of oxidized ZS2 at 19001C for 1 h. Thicknesses of the scale of ZS1, ZS2, and ZS3 after oxidation at 19001C in air for 1 h are 1100, 600, and 800 mm, respectively. The thickness of the oxide scale of ZS2 is  lower than ZS1 and ZS3. Interestingly, ZS2 is more resistant to  oxidation than ZS3, which is contrary to the results at lower temperature.20 The scale for ZS2 remains coherent and attached  to the unreacted material despite the formation of voids in the  SiC-depleted region. Furthermore,  the microstructure of  the  subscale for ZS2 was remarkably different from ZS3. The sub scale of ZS2 shows an oriented growth. The crystalline zirconia  exhibited a  columnar  structure  in the oxidized ZS2  sample  (Fig. 6(c)), whereas this phenomenon was not observed in ZS3  after oxidation in the same condition. The oriented growth of  the scale is most likely due to the evolution of the gaseous by products, which promoted growth of zirconia in the direction  parallel to the discharge of the gaseous products. The preferential active oxidation of SiC at 19001C did not provide enough  SiO2 for passive oxidation of ZrB2-SiC composites. In addition, high SiC content would promote the formation of a large SiC depleted zone and the transport of the oxide products of ZrB2, which are detrimental to the oxidation of ZrB2-SiC. Figure 7 shows the enlarged view of the cross sections for ZS1, ZS2, and ZS3 after oxidation at 19001C. Cracks and spallation were de tected at the interface between unreacted material and the oxide  scale for ZS1. A dense adherent oxide scale composed of ZrO2 and SiO2 was formed for ZS2, whereas a loose porous structure was generated for ZS3. Moreover, cracks and spallation beneath  the surface layer were also observed. The images of  the cross  sections further revealed a strong degradation of ZS1 and ZS3  relative to ZS2. It should be noted that  the failure modes and  mechanisms of  the ZrB2-SiC composites with low (e.g., ZS1) and high (e.g., ZS3) SiC content are different.  Thermodynamically, both ZrB2 and SiC should be cantly oxidized in the present case. The oxide products of SiC  signifi would either escape as gas phases or be transported to the sur face  layer  through liquid convection. Consequently, SiC has  been entirely removed from the partially depleted zirconium  diboride layer. The oxide products of ZrB2 are ZrO2 and B2O3. In fact, the last compound has an unusually low melting point (4501C) and a high vapor pressure. Therefore, at high-temper ature B2O3 quickly vaporizes. Most likely, for the transport of ZrO2, then the formed ZrO2 would not adhere to the unreacted ZrB2 at 19001C, which is significantly different from the oxidation behavior of the ZrB2-based UHTC at lower temperature. The formation and migration of the liquid  if there is a channel  ZrO2 oxides at this phenomenon must be taken into account when the oxidation temperature is very high (e.g., 19001C). ZrO2 can form liquid oxides with SiO2 at high temperature, and this can be quantitatively analyzed from the phase diagram of the binary ZrO2-SiO2 system.21 Therefore, ZrO2 can be transported to the surface layer by convection in the liquid silica, as well as disso lower temperature can be neglected. However,  lution in and recrystallization from the liquid silica. In fact, when the system is above the boiling point of B2O3 (18601C), oxidation resistance of ZrB2 decreases remarkably, and the formed oxidation products will not be adhered to the unaltered  ZrB2. The consumption of SiC would leave behind channels for the transport of the formed ZrO2. Moreover, the formation of high-pressure gaseous products would also accelerate the trans port of ZrO2, resulting in the generation of a discrete structure as shown in Figs. 3 and 4. However, this phenomenon was not  observed in ZS2 (Figs. 6(a) and (d)). The structure in the SiC depleted region was also continuous for ZS2, which is attributed  to the difference of SiC content.  300 µm  (a)  50 µm  (b)  Silica-rich layer  Zirconia-rich layer  SiC-depleted layer  Unreacted material  Fig. 3. Cross-sectional scanning electron micrographs of ZS3 oxidized at 19001C in air for 1 h (a), and a detail of the SiC-depleted layer (b).  3330  Journal of the American Ceramic Society—Han et al.  Vol. 91, No. 10  \\x0c', 'October 2008  High-Temperature Oxidation at 19001C of ZrB2-xSiC  3331  (a)  (c)  ZrO2  (b)  (d)  ZrB2  SiO2  200 µm  ZrB2  10 µm  10 µm  ZrO2  50 µm  Fig. 4.  Scanning electron micrographs of the subside scale of ZS3 after successfully removing the outside scale.  (a)  (b)  300 µm  20 µm  Fig. 5. Cross-sectional scanning electron micrographs of ZS1 oxidized at 19001C in air for 1 h.  (3)  Effect of SiC Content on the Oxidation Resistance  Theoretically, ZrB2-SiC consists of a ZrB2 matrix and dispersed SiC particles. No significant solid solution is expected between ZrB2 and SiC, which are stable when in contact with each other.19 A schematic oxidation model of ZrB2 containing low (e.g., 10 vol%), medium (e.g., 20 vol%), and high (e.g., 30 vol%) SiC  content is shown in Fig. 8. This paper represents the first attempt  to use a model to interpret the effect of SiC content on the ox idation resistance of ZrB2-SiC UHTC composites. Thermodynamically, both ZrB2 and SiC should be oxidized when exposed to air. The dominant chemical process at 19001C for ZrB2 is the oxidation according to reaction (1), which produces an oxide  layer that mainly consists of ZrO2. The expected main reactions for SiC are reactions (2) and (3). At 19001C, reaction (3) is fa vored. In fact, SiC formed a network in both ZS2 and ZS3, which  was interconnected in three dimensions, although it was discon tinuous in two dimensions. The degree of SiC interconnectivity in  the matrix increases with increasing SiC content. Kinetically, the  oxidation of SiC is more rapid than the oxidation of ZrB2 at this temperature, as evidenced by ZrB2 inclusions present in the scale.  ZrB2 ðsÞ þ 5 2 O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðgÞ  SiCðsÞ þ 3 2 O2 ðgÞ ! SiO2 ðl Þ þ COðgÞ  SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  (1)  (2)  (3)  SiC exhibits preferential oxidation in the present  system of  leading to the formation of the layer depleted in SiC  ZrB2-SiC, grains. At  the same time, ZrB2 grains that are in contact with depleted SiC grains are directly exposed to oxygen. Thus, the  oxidation occurs from the boundary between ZrB2 and SiC to inner ZrB2 grains, which leads to the development of a structure similar to that shown schematically in Fig. 8. It should be noted  that SiC particles  are  assumed to be of  the  same  size  and  homogenously  dispersed within the matrix. Apparently,  the  degree of SiC interconnectivity in the matrix increases with in creasing SiC content. A high degree of SiC interconnectivity  \\x0c', '3332  Journal of the American Ceramic Society—Han et al.  Vol. 91, No. 10  (a)  (c)  (b)  (d)  300 µm  50 µm  50 µm  50 µm  Fig. 6.  Cross-sectional scanning electron micrographs of ZS2 oxidized at 19001C in air for 1 h (a); (b), (c), and (d) are the details of the silica-rich layer,  zirconia-rich layer, and SiC-depleted layer, respectively.  causes  the  rapid active oxidation of SiC. Consequently,  the  thickness of  the SiC-depleted layer  for ZS3  is much higher  Based on the relative oxidation rates of ZrB2 and SiC, SiC should be consumed first through the rapid active oxidation. On  than that  for ZS2 under  the  same  conditions,  as  shown in  the one hand, the SiC consumption left behind channels for the  Figs. 3 and 6.  transport of the formed ZrO2. On the other hand, the SiC con (a)  Cracks / spallation   (b)  (c)  500 µm  Cracks / spallation   300 µm  300 µm  Fig. 7.  Enlarged view of the cross sections for (a) ZS1, (b) ZS2, and (c) ZS3 after oxidation at 19001C in air for 1 h.  \\x0c', 'sumption left behind space for the volume expansion resulting  from the oxidation of ZrB2. Cracks or spallation would occur when the vacant space cannot satisfy the volume expansion  upon the conversion of ZrB2 to ZrO2 (Fig. 7(a)). The boiling of B2O3 and continuous evolution of CO and SiO would make the formed ZrO2 phase nonadherent. ZrO2 is not stationary in the present case and will be moved to the external oxide scale  through the  initially formed,  connected pores. Apparently,  if  the oxidized regimes of  the ZrB2 matrix within the SiC-depleted region and their oxidation products  are  interconnected  are transported to the external oxide scale, the solid skeleton will  not  interconnect with itself when the holes meet  together  (see  Fig. 8(c)). Consequently,  the scale for  this material  is visually  nonadherent and some oxide is spallated (Figs. 3 and 8(c)). On  the other hand, the rupture of the scale may not occur (Fig. 6)  when the holes in the ZrB2 matrix are not interconnected in the SiC-depleted region, as shown in Fig. 8(b). Whether the oxidized  regimes of the ZrB2 matrix in the SiC-depleted region are interconnected or not mostly depends on the content of SiC in the  matrix (see Figs. 7 and 8). A SiC-depleted layer would not be  formed when the SiC content  is too low to form interconnec tivity (see Fig. 8(a)). The average thickness of ZrB2 matrix between two adjacent SiC particles for high SiC content (ZS3) is  lower  than that of medium SiC content  (ZS2), as  shown in  Fig. 8. Therefore,  the oxidized regimes of  the ZrB2 matrix in high SiC content (ZS3) are more liable to meet together in com parison with those of low SiC content (ZS2). ZS2 exhibited superior oxidation resistance at 19001C in air  for 1 h. However, the oxidation behavior of ZS1 and ZS3 under  the same conditions demonstrated the unsuitability of this ma terial for ultrahigh-temperature applications. Comparison of the  oxidation results between ZS1, ZS2, and ZS3 indicates that SiC  content significantly affected the oxide structure and performance of ZrB2-SiC at 19001C. Neither high (e.g., 30 vol% SiC) nor low (e.g., 10 vol% SiC) content of SiC is appropriate for  ultrahigh-temperature applications. With the assumption that  oxidation products have theoretical density, 1 unit volume of  ZrB2 upon oxidation produces 1.19 unit volumes of solid oxide (ZrO2). For ZrB2 containing 10% SiC composite, the volume increase of the solid oxide during oxidation is 7%, leading to the  cracks or  spallation (B2O3 and SiO2 were not count). In addition, the amount of the silica is lower than the  taken into ac high-SiC-content materials  (i.e., ZS2 and ZS3). Obviously, a  small amount of silica glass cannot provide an effective barrier  for oxygen ingress. Similarly, ZrB2 containing 30% SiC composite is likely to crack and spall due to the failure of the SiC depleted layer. However, the resulting structure of both the SiC depleted layer and the ZrO2-rich layer crack and spall and remains more adherent. With respect  in ZS2 is  less likely to  to  ZS2, the fraction of the formed solid is 95%, which will not lead  to the cracks or spallation as a result of volume expansion dur ing oxidation. The other fraction will be filled with liquid silica  glass, which provides an effective obstacle against  the inward  diffusion of oxygen along short-circuit paths (i.e., residual po rosity, cracks, and grain boundary) and it is helpful in forming a  dense coherent oxide scale at the same time. Moreover, cracks  and spallation would not occur in the SiC-depleted layer under  the present condition, which was  significantly affected by the  SiC content. How to obtain a dense adherent oxide scale be comes a major issue for  the applications of  these materials at  ultrahigh temperature, especially 19001C. According to the previous  for  temperatures  above  study, high SiC content  is  beneficial to the oxidation resistance of the material below 18001C, whereas it is detrimental at higher temperature. There fore, we should optimize the SiC content and microstructure to  meet the requirements of the practical application.  IV.  Conclusions  Dense ZrB2-SiC UHTC composites were prepared by hot pressing. ZrB2 containing 20 vol% SiC particulates exhibited the highest oxidation resistance at 19001C in air  compared with  ZrB2-10 vol% SiC and ZrB2-30 vol% SiC composites. ZrO2 is not stationary at high temperature and can be transported di rectly or by liquid convection to the surface layer, which is sig nificantly different from that at lower temperature. The resulting  structure of both the SiC-depleted layer and the ZrO2-rich layer in ZrB2-20 vol% SiC is less likely to crack and spall and remains more adherent. SiC content significantly affected the formation  of the oxide structure and the performance of the UHTC at this  temperature. These UHTC materials with low or high SiC con tent are  inappropriate  for ultrahigh-temperature applications.  The SiC content and microstructure in the ZrB2 matrix should be optimized according to the requirements of the practical  application.  References  1K. Upadhya, J. M. Yang, and W. P. Hoffman, ‘‘Materials for Ultrahigh Tem perature Structural Applications,’’ Am. Ceram. Soc. Bull., 76 [12] 51-6 (1997).  Unreacted material  ZrB2  SiC  Zirconia-rich layer  Silica-rich layer  O2(g)  O2(g)  CO (g), SiO(g)  B2O3(g)  B2O3(g)  ZrO2  ZrO2  ZrO2  ZrB2  SiC  Zirconia-rich layer  Silica-rich layer   ZrO2  Unreacted material  SiC-depleted layer   CO (g), SiO(g)  O2(g)  ZrO2  ZrB2  SiC  SiC-depleted layer   Zirconia-rich layer   Silica-rich layer  ZrO2  Unreacted material  CO (g), SiO(g)  B2O3(g)  (a)  (b)  (c)  Fig. 8. Oxidation models of the present SiC reinforced ZrB2 matrix ultrahigh-temperature ceramics at 19001C: (a) low SiC content (ZS1), (b)  medium SiC content (ZS2), and (c) high SiC content (ZS3); white color  represents holes.  October 2008  High-Temperature Oxidation at 19001C of ZrB2-xSiC  3333  \\x0c', '3334  Journal of the American Ceramic Society—Han et al.  Vol. 91, No. 10  2A. K. Kuriakose and J. L. Margrave,  ‘‘The Oxidation Kinetics of Zirconium  12X. H. Zhang, P. Hu, J. C. Han, L. Xu, and S. H. Meng,  ‘‘The Addition of  Diboride and Zirconium Carbide at High Temperature,’’ J. Electrochem. Soc., 111  Lanthanum Hexaboride to Zirconium Diboride for Improved Oxidation Resis [7] 827-31 (1964). 3S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem,  tance,’’ Scr. Mater., 57 [11] 1036-9 (2007). 13D. M. Van Wie, D. G. Jr. Drewry, D. E. King, and C. M. Hudson,  ‘‘The  ‘‘Evaluation  of Ultra-High Temperature Ceramics  for Aeropropulsion Use,’’  Hypersonic Environment: Required Operating Conditions  and Design Chal J. Euro. Ceram. Soc., 22 [14, 15] 2757-67 (2002). 4M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties and Arc Jet Oxidation of Hafnium Diboride/Silicon Car bide Ultra High Temperature Ceramic,’’ J. Mater. Sci., 39 [19] 5925-37 (2004). 5F. Monteverde and A. Bellosi,  ‘‘The Resistance to Oxidation of an HfB2-SiC Composite,’’ J. Eur. Ceram. Soc., 25 [7] 1025-31 (2005). 6A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Oxidation of  ZrB2-SiC Ceramics Under Atmospheric and Reentry Conditions,’’ Refract. Appl. Trans., 1 [2] 1-8 (2005). 7W. G. Fahrenholtz, G. E. Hilmas, A. L. Chamberlain, and J. W. Zimmermann,  ‘‘Processing and Characterization of ZrB2-Based Ultra-High Temperature Monolithic and Fibrous Monolithic Ceramics,’’ J. Mater. Sci., 39, 5951-7 (2004). 8F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 9F. Monteverde, ‘‘The Thermal Stability in Air of Hot-pressed Diboride Matrix  in Dry  Composites for Uses at Ultra-High Temperatures,’’ Corrosion Sci., 47 [8] 2020-33  (2005). 10F. Monteverde and A. Bellosi,  ‘‘Microstructure and Properties of an HfB2- SiC Composite for Ultra High Temperature Applications,’’ Adv. Eng. Mater., 6 [5]  331-6 (2004). 11E. J. Opila and M. C. Halbig,  Proc., 22 [3] 221-8 (2001).  ‘‘Oxidation of ZrB2-SiC,’’ Ceram. Eng. Sci.  lenges,’’ J. Mater. Sci., 39 [19] 5915-24 (2004). 14M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurfaces: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 15A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  ‘‘Characterization  of Zirconium Diboride-Molybdenum Disilicide Ceramics,’’  Ceram. Trans., 153, 299-308 (2004). 16E. J. Opila, S. R. Levine, and J. Lorincz, ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39,  5969-77 (2004). 17S. R. Levine and E. J. Opila,  ‘‘Tantalum Addition to Zirconium Diboride for  Improved Oxidation Resistance,’’ NASA/TM—2003-212483. 18W. G. Fahrenholtz,  ‘‘Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depletion Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 19A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Evolution of Structure  during the Oxidation of Zirconium Diboride-Silicon Carbide 15001C,’’ J. Eur. Ceram. Soc., 27, 2495-501 (2007). 20J. C. Han, P. Hu, X. H. Zhang, and S. H. Meng, ‘‘Oxidation Behavior of 18001C,’’ Scr. Mater.,  Zirconium Diboride-Silicon Carbide  in Air Up  825-8  to  at  57  [9]  (2007). 21W. C. Butterman and W. R. Foster,  ‘‘Zircon Stability and the ZrO2-SiO2 Phase Diagram,’’ Am. Mineral., 52, 880-5 (1967).  &  \\x0c']"
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  "PDF": "HIGH-TEMPERATURE OXIDATION OF ZIRCONIUM CARBIDE AT LOW OXYGEN-PRESSURE.pdf",
  "Text": "['Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   H i g h T e m p e r a t u r e   O x i d a t i o n   I V .   Z i r c o n i u m   a n d   H a f n i u m   C a r b i d e s   J o a n   B .   B e r k o w i t z M a t t u c k   A r t h u r   D .   L i t t l e ,   I n c . ,   C a m b r i d g e ,   M a s s a c h u s e t t s   A B S T R A C T   T h e   o x i d a t i o n   o f   Z r C   w a s   s t u d i e d   a t   t e m p e r a t u r e s   o f   i 1 3 0 ~ 1 7 6   a n d   o x y g e n   p a r t i a l   p r e s s u r e s   a r o u n d   3 . 9   a n d   2 0   T o r r .   T h e   r a t e   o f   o x i d a t i o n   w a s   m o n i t o r e d   w i t h   a   t h e r m a l   c o n d u c t i v i t y   c e l l .   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T h e   o x i d a t i o n   o f   z i r c o n i u m   c a r b i d e   w a s   s t u d i e d   b y   M a r g r a v e   a n d   K u r i a k o s e   ( 1 )   a t   t e m p e r a t u r e s   o f   5 5 0   ~   6 5 0 ~   i n   o x y g e n   a t   1   a t m .   T h e   o x i d a t i o n   w a s   f o u n d   t o   b e   l i n e a r ,   w i t h   a n   a c t i v a t i o n   e n e r g y   o f   1 6 . 7   _   1 . 7   k c a l / m o l e .   B a r t l e t t ,   W a d s w o r t h ,   a n d   C u t l e r   ( 2 )   s t u d i e d   t h e   w e i g h t   g a i n   o f   s i z e d   p o w d e r s   o f   z i r c o n i u m   c a r b i d e   i n   a i r ,   o x y g e n ,   a n d   o x y g e n h e l i u m   m i x t u r e s   a t   t e m p e r a   t u r e s   o f   4 5 0 ~ 1 7 6   a n d   o x y g e n   p r e s s u r e s   o f   6 . 5   x   1 0 ~   t o   1   a t m .   S t o i c h i o m e t r i c   o x i d a t i o n   o f   Z r C   t o   Z r O 2   a n d   g a s e o u s   o x i d e s   o f   c a r b o n   w a s   a s s u m e d .   T h e   d a t a   w e r e   i n t e r p r e t e d   o n   t h e   b a s i s   o f   t w o   p a r a l l e l   i n d e   p e n d e n t   p r o c e s s e s :   a   p a r a b o l i c   d i f f u s i o n   r e a c t i o n   i n   v o l v i n g   t h e   p a r t i a l   r e p l a c e m e n t   o f   i n t e r s t i t i a l   c a r b o n   i n   t h e   Z r C   l a t t i c e   w i t h   o x y g e n ,   a n d   a   l i n e a r   s u r f a c e   r e a c t i o n   o c c u r r i n g   a t   t h e   Z r C Z r O 2   p h a s e   b o u n d a r y .   B o t h   p r o c e s s e s   o c c u r   s i m u l t a n e o u s l y ,   w i t h   t h e   d i f f u s i o n   r e a c t i o n   p r e d o m i n a t i n g   a t   s h o r t   t i m e s   a n d   t h e   s u r f a c e   r e a c t i o n   b e c o m i n g   r a t e   c o n t r o l l i n g   a s   o x i d a t i o n   p r o   c e e d s .   T h e   a c t i v a t i o n   e n e r g i e s   w e r e   c a l c u l a t e d   a s   5 3   k c a l / m o l e   f o r   t h e   d i f f u s i o n   p r o c e s s   a n d   4 6   k c a l / m o l e   f o r   t h e   s u r f a c e   r e a c t i o n .   W a t e r   v a p o r ,   i n   t h e   p r e s e n c e   o f   o x y g e n ,   w a s   f o u n d   t o   a c c e l e r a t e   t h e   r a t e   o f   t h e   s u r   f a c e   r e a c t i o n   w h i l e   l e a v i n g   i t s   a c t i v a t i o n   e n e r g y   u n   c h a n g e d .   T h e   s o l i d   o x i d a t i o n   p r o d u c t   w a s   f o u n d   t o   b e   c u b i c   Z r O 2 ,   a   p h a s e   n o r m a l l y   t h o u g h t   t o   b e   u n s t a b l e ,   b u t   w h i c h   m i g h t   b e   s t a b i l i z e d   b y   s m a l l   a m o u n t s   o f   c a r b o n .   W a t t ,   C o c k e t t ,   a n d   H a l l   ( 3 )   m a d e   a   s i n g l e   w e i g h t   c h a n g e   m e a s u r e m e n t   o f   4 9 . 8   m g / c m ~   o n   a   s o l i d   s a m p l e   o f   Z r C   o f   d e n s i t y   6 . 2 0   g / c c   a n d   4 . 8 %   p o r o s i t y ,   e x   p o s e d   t o   a   s t r e a m   o f   d r y   a i r   f l o w i n g   a t   5 . 3   c m / s e c ,   f o r   3 0   m i n   a t   8 0 0 ~   T h e   p r e s e n t   s t u d y   w a s   u n d e r t a k e n   t o   i n v e s t i g a t e   t h e   o x i d a t i o n   o f   Z r C   a n d   t h e   c h e m i c a l l y   r e l a t e d   H f C   a t   t e m p e r a t u r e s   a b o v e   9 0 0 ~   E x p e r i m e n t a l   M e t h o d   C y l i n d r i c a l   p e l l e t s   o f   Z r C   w e r e   c u t   f r o m   z o n e   r e   f i n e d   b a r s   p r e p a r e d   a s   d e s c r i b e d   b y   W e s t r u m   a n d   F e i c k   ( 4 ) .   T h e   f a b r i c a t e d   b a r s   c o n t a i n e d   1 1 . 2   w / o   c a r b o n .   H a f n i u m   c a r b i d e   p o w d e r   w a s   p r e p a r e d   b y   t h e   C a r b o r u n d u m   C o m p a n y   f r o m   h i g h p u r i t y   H f O ~   s u p   p l i e d   b y   W a h   C h a n g   C o r p o r a t i o n .   T h e   H f C   p o w d e r   w a s   s i n t e r e d   i n t o   b a r s   a n d   a r c m e l t e d   o n   a   w a t e r   c o o l e d   c o p p e r   h e a r t h   u s i n g   a   w a t e r c o o l e d   t u n g s t e n   e l e c t r o d e .   I n   o r d e r   t o   m i n i m i z e   l o s s   o f   c a r b o n   d u r i n g   m e l t i n g ,   t h e   o p e r a t i o n   w a s   c o n d u c t e d   i n   a n   a t m o s   p h e r e   o f   a r g o n   c o n t a i n i n g   3 . 1 4 %   o f   e t h y l e n e   a n d   1 1 . 4 %   o f   h y d r o g e n .   T h e   r e s u l t i n g   m a t e r i a l   w a s   c a r b o n   d e   f i c i e n t ,   c o r r e s p o n d i n g   t o   a   c o m p o s i t i o n   H f C 0 . 9 5 2   ( 5 ) .   T h e   a n t i c i p a t e d   p r o d u c t s   o f   t h e   o x i d a t i o n   o f   Z r C   a n d   H f C   w e r e   t h e   p e r m a n e n t   g a s e s ,   C O ( g )   a n d   C O 2 ( g ) ,   i n   a d d i t i o n   t o   t h e   r e f r a c t o r y   m e t a l   o x i d e s .   D u e   t o   t h e   e v o l u t i o n   o f   p e r m a n e n t   g a s e s ,   t h e   t h e r m a l   c o n d u c   t i v i t y   m e t h o d   d e s c r i b e d   i n   p r e v i o u s   p u b l i c a t i o n s   ( 6 8 )   h a d   t o   b e   m o d i f i e d   t o   s t u d y   t h e   o x i d a t i o n .   A   k n o w n   m i x t u r e   o f   h e l i u m   a n d   o x y g e n   w a s   p a s s e d   t h r o u g h   t h e   r e f e r e n c e   s i d e   o f   a   t h e r m a l   c o n d u c t i v i t y   c e l l   ( 6 )   a n d   o v e r   a n   i n d u c t i v e l y   h e a t e d   c a r b i d e   p e l l e t   s u p p o r t e d   o n   T h O 2   f i n g e r s   b y   t h r e e   p o i n t   c o n t a c t .   A   p o r t i o n   o f   t h e   o x y g e n   i n   t h e   g a s   s t r e a m   r e a c t e d   w i t h   t h e   p e l l e t   t o   p r o d u c e   o x i d e s   o f   t h e   m e t a l   a n d   c a r b o n .   T h e   e f f l u e n t   g a s   w a s   t h e r e f o r e   d e p l e t e d   i n   o x y g e n ,   b u t   e n r i c h e d   i n   C O ( g )   a n d   C O 2 ( g ) .   T h e   l a t t e r   w a s   r e m o v e d   b y   p a s s   a g e   t h r o u g h   a   w e i g h e d   A s c a r i t e   b u l b ,   a n d   t h e   r e m a i n   i n g   m i x t u r e   o f   C O ( g ) ,   O 2 ( g ) ,   a n d   H e   e n t e r e d   t h e   s a m p l i n g   s i d e   o f   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   ( 6 ) .   F i n a l l y ,   t h e   C O ( g )   w a s   o x i d i z e d   t o   C O 2 ( g )   o v e r   c o p   p e r   o x i d e   p o w d e r   a t   5 0 0 ~   a n d   t h e   C O s ( g )   p r o d u c e d   w a s   c o l l e c t e d   i n   a   s e c o n d   w e i g h e d   A s c a r i t e   b u l b .   T h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   i n   t h i s   c a s e   w a s   t h e r e f o r e   n o t   d i r e c t l y   r e l a t e d   t o   t h e   r a t e   o f   o x y   g e n   c o n s u m p t i o n   a s   i n   p r e v i o u s l y   s t u d i e d   s y s t e m s   ( 6 8 ) ,   b u t   i n s t e a d   r e f l e c t e d   t h e   d i f f e r e n c e   b e t w e e n   t h e   r a t e   o f   t o t a l   o x y g e n   c o n s u m p t i o n   a n d   t h e   r a t e   o f   f o r m a t i o n   o f   C O ( g ) .   T h e   c a l i b r a t i o n   c o n s t a n t ,   r e   l a t i n g   t h e   e l e c t r i c a l   s i g n a l   t o   t h e   d i f f e r e n c e   i n   g a s   p r e s s u r e   o n   t h e   t w o   s i d e s   o f   t h e   c e l l   i s   t h e   s a m e   f o r   b o t h   C O   a n d   O ~   i n   H e .   H e n c e ,   e v o l u t i o n   o f   C O   d e   p r e s s e s   t h e   s i g n a l ,   a s   c o m p a r e d   t o   s i m p l e   r e m o v a l   o f   o x y g e n .   R e s u l t s   Z i r c o n i u m   c a r b i d e . O x i d a t ~ e n   k i n e t i c s . T h e   e x   p e r i m e n t a l   d a t a   a r e   s u m m a r i z e d   i n   T a b l e   I .   T h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   a p p a r a t u s   w a s   c o n s t a n t   w i t h   t i m e   i n   e v e r y   e x p e r i m e n t .   H o w e v e r ,   o n l y   i n   t h e   c a s e   o f   e x p e r i m e n t   X I I 1   d i d   t h i s   r e f l e c t   t r u e   c h e m i   c a l l y   c o n t r o l l e d   l i n e a r   o x i d a t i o n   k i n e t i c s .   I n   t h e   o t h e r   e x p e r i m e n t s   m o r e   t h a n   9 0 %   o f   t h e   o x y g e n   p a s s e d   o v e r   t h e   c a r b i d e   p e l l e t   r e a c t e d   w i t h   i t ,   a n d   t h e   c o n t r o l l i n g   p r o c e s s   w a s   t h e r e f o r e   p r o b a b l y   t h e   r a t e   o f   a r r i v a l   o f   o x y g e n   g a s   a t   t h e   s a m p l e   s u r f a c e .   A t   h i g h e r   p r e s s u r e s   a n d / o r   h i g h e r   g a s   f l o w   r a t e s ,   a   g r e a t e r   p r o p o r t i o n   o f   t h e   c a r b i d e   w o u l d   h a v e   b e e n   c o n v e r t e d   t o   o x i d e s .   I n   T a b l e   I ,   t h e   \" i n i t i a l \"   w e i g h t s   w e r e   t a k e n   a f t e r   d e g a s s i n g   a t   2 2 0 0 ~   i n   p u r e   h e l i u m   u n t i l   t h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   i n d i c a t e d   t h a t   n o   p e r m a n e n t   g a s e s   w e r e   b e i n g   e v o l v e d .   T h e   s u r f a c e   a r e a s   w e r e   c a l c u l a t e d   f r o m   m i c r o m e t e r   m e a s u r e m e n t s   o f   t h e   h e i g h t   a n d   d i a m e t e r   o f   t h e   c y l i n d r i c a l   p e l l e t s .   T e m p e r a t u r e s   w e r e   m e a s u r e d   o p t i c a l l y   a n d   c o r r e c t e d   f o r   a n   e m i s s i v i t y   o f   0 . 7 ,   d e t e r m i n e d   b y   c o m p a r i n g   1 0 3 0   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   V o l .   1 1 4 ,   N o .   1 0   H I G H T E M P E R A T U R E   O X I D A T I O N   T a b l e   I .   S u m m a r y   o f   e x p e r i m e n t a l   d a t a   o n   Z r C   1 0 3 1   G e o m e t r i c   I n i t i a l   s u r f .   a r e a ,   E x p t .   w e i g h t ,   g   c m   2   T e m p ,   * I 4 [   C O z   O x y g e n   E x p o s u r e   N e t   w e i g h t   C O   f o r m e d   C   Z r   p r e s s u r e ,   t i m e ,   c h a n g e   f o r m e d   W c o   ( g )   c o n s u m e d   c o n s u m e d   T o r r   r a i n   W o ,   ( g )   W o o   ( g )   W e   ( g )   W z r   ( g )   Z r / C   X I I 8   0 . 4 4 6   1 . 1 0 3   1 1 3 0   X I I 5   0 . 5 6 4 5   1 . 3 9 4   1 2 6 0   X I I 3   0 . 6 9 6 3   1 . 4 9 3   1 5 6 0   X 3 1   0 . 5 8 6 4   1 . 3 1 0   1 8 6 0   X 2 9   0 . 7 0 7 3   1 . 4 5 0   1 9 4 0   X I I I   0 . 7 0 5 1   1 . 5 0 0   1 9 7 0   X 2 7   0 . 6 3 6 1   1 . 2 5 2   2 0 7 0   X 1 6   0 . 6 8 4 0   1 . 3 8 6   2 0 7 0   V I I 8   0 . 6 6 4 3   1 . 5 2 8   2 1 0 0   X 2 5   0 . 6 8 1 2   1 . 3 7 1   2 1 6 5   2 2 . 9   5 1     0 . 0 0 0 1   0 . 0 6 1 1         2 0 . 4   1 2 8       0 . 0 7 0 9     ~     2 1 . 2   6 2     0 . 0 2 8 9   0 . 0 3 9 9       9 . 1   1 2 9   0 . 0 4 0 2   0 . 0 4 3 8   0 . 0 2 2 0   0 . 0 2 4 8   0 . 1 8 5 2   8 . 1   1 2 0   0 . 0 3 6 6   0 . 0 4 0 5   0 . 0 1 4 5   0 . 0 2 1 4   0 . 1 6 5 2   7 . 7 2   2 5 . 9   1 1 2   0 . 0 8 5 1   0 . 0 9 6 8   0 . 0 3 5 6   0 . 0 5 1 0   0 . 3 8 8   7 . 6 0   8 . 5   1 2 4   0 . 0 3 8 9   0 . 0 4 6 4   0 . 0 1 7 5   0 . 0 2 3 6   0 . 1 7 8 0   7 . 5 5   6 . 5   1 1 9   0 . 0 4 0 2   0 . 0 5 4 0   0 . 0 0 7 3   0 . 0 2 5 2   0 . 1 8 6 2   7 . 3 9   3 . 0   1 8 0   0 . 0 2 0 7   0 . 0 2 4 2   0 . 0 0 7 9   0 . 0 1 2 6   0 . 0 9 4 9   7 . 5 4   8 . 9   1 2 0   0 . 0 3 5 6   0 . 0 5 2 3   0 . 0 0 7 2   0 . 0 2 4 4   0 . 1 7 1 0   7 . 0 1   s u r f a c e   t e m p e r a t u r e s   w i t h   u l t r a s o n i c a l l y   d r i l l e d   b l a c k   b o d y   c a v i t y   t e m p e r a t u r e s   u n d e r   o x i d i z i n g   c o n d i t i o n s .   O x y g e n   p a r t i a l   p r e s s u r e s   i n   h e l i u m   a r e   g i v e n ;   t h e   t o t a l   p r e s s u r e   w a s   c l o s e   t o   1   a t m   i n   e v e r y   e x p e r i m e n t .   T h e   c a r r i e r   g a s   f l o w   r a t e   w a s   5 8 . 6   c c / m i n   i n   e v e r Y   e x   p e r i m e n t ,   c o r r e s p o n d i n g   t o   a   l i n e a r   f l o w   v e l o c i t y   i n   t h e   n e i g h b o r h o o d   o f   t h e   s a m p l e   o f   1 . 9   c m / s e c ,   e x c e p t   V I I 8   w h e r e   a   f l o w   o f   5 1 . 5   c c / m i n   w a s   u s e d .   F r o m   t h e   m e a s u r e d   n e t   w e i g h t   c h a n g e   o f   t h e   c a r   b i d e   o n   o x i d a t i o n ,   a n d   t h e   o b s e r v e d   w e i g h t   c h a n g e s   W c o 2   a n d   W c o   i n   t h e   A s c a r i t e   b u l b s ,   t h e   t o t a l   a m o u n t s   o f   c a r b o n   a n d   z i r c o n i u m   c o n s u m e d   w e r e   r e a d i l y   c a l   c u l a t e d ,   o n   t h e   a s s u m p t i o n   t h a t   t h e   o n l y   o x i d a t i o n   p r o d u c t s   w e r e   C O 2 ( g ) ,   C O ( g ) ,   a n d   Z r O ~ ( s ) .   T h e   f o r m a t i o n   o f   Z r O C   ( s )   c a n n o t   b e   p r e c l u d e d ,   b u t   s i n c e   i t   i s   i s o s t r u c t u r a l   w i t h   Z r C ,   n o   p o s i t i v e   e v i d e n c e   w a s   o b   t a i n e d   f o r   i t s   p r e s e n c e .   T h e   t o t a l   w e i g h t   o f   c a r b o n   c o n s u m e d ,   W e ,   i s   g i v e n   b y   [ c ]   [ c ]   W e   =     (cid:12) 9   W c o 2   +   ~   \" W c o   [ 1 ]   [ C O ~ ]   [ C O ]   w h e r e   t h e   s y m b o l s   i n   b r a c k e t s   r e p r e s e n t   m o l e c u l a r   w e i g h t s .   T h e   t o t a l   w e i g h t   o f   z i r c o n i u m ,   W z r ,   t h a t   h a s   b e e n   c o n v e r t e d   t o   o x i d e   i s   c a l c u l a t e d   f r o m   t h e   m e a   s u r e d   w e i g h t   c h a n g e ,   W o ,   a n d   t h e   d e r i v e d   c a r b o n   c o n   s u m p t i o n   [ Z r ]   W z r   =   ~   ( W o   +   W c )   [ 2 ]   2 [ 0 ]   T h e   r a t i o   o f   t h e   n u m b e r   o f   g r a m s   o f   z i r c o n i u m   c o n   s u m e d   t o   t h e   n u m b e r   o f   g r a m s   o f   c a r b o n   c o n s u m e d   d u r i n g   o x i d a t i o n   i s   s h o w n   i n   T a b l e   I   t o   h a v e   a n   a p   p r o x i m a t e l y   c o n s t a n t   v a l u e   o f   7 . 5   (cid:127)   0 . 2 .   S i n c e   t h e   c o r r e s p o n d i n g   r a t i o   i n   t h e   Z r C   s t a r t i n g   m a t e r i a l   i s   7 . 6 ,   i t   w o u l d   a p p e a r   t h a t   t h e   o x i d a t i o n   o f   Z r C   i s   s t o i   c h i o m e t r i c .   T h a t   i s ,   f o r   e a c h   z i r c o n i u m   a t o m   c o n v e r t e d   t o   o x i d e ,   a   s i n g l e   c a r b o n   a t o m   i s   a l s o   c o n v e r t e d   t o   o x i d e .   S t r u c t u r a l   c h a n g e s   d u r i n g   o x i d a t i o n . T h e   r e a s o n   t h a t   w e i g h t   c h a n g e   d a t a   w e r e   n o t   g i v e n   f o r   p e l l e t s   X I I 8 ,   X I I 5 ,   a n d   X I I 3   i s   t h a t   a t   t h e s e   r e l a t i v e l y   l o w   t e m p e r a t u r e s   t h e   p e l l e t s   w e r e   b r o k e n   a p a r t   b y   t h e   o x i d a t i o n   p r o c e s s .   A t   t h e   e n d   o f   e a c h   e x p e r i m e n t ,   t h e   g r a i n   b o u n d a r i e s   o f   t h e   Z r C   w e r e   s e e n   t o   b e   o u t l i n e d   b y   a   w h i t e   m a t e r i a l ,   i d e n t i f i e d   b y   r o o m   t e m p e r a t u r e   x r a y   a s   m o n o c l i n i c   Z r O 2 .   T h e   g r o w t h   o f   t h e   o x i d e   i n   p r e e x i s t i n g   c r a c k s   a n d   g r a i n   b o u n d a r i e s   o f   Z r C   u n   d o u b t e d l y   c r e a t e s   e n o u g h   s t r e s s   t o   f r a c t u r e   t h e   c a r   b i d e .   A t   h i g h e r   t e m p e r a t u r e s ,   a s   d i s c u s s e d   b e l o w ,   s i g   n i f i c a n t   g r a i n   b o u n d a r y   o x i d a t i o n   w a s   o b s e r v e d ,   b u t   s t r e s s e s   a r e   a p p a r e n t l y   s u f f i c i e n t l y   r e l i e v e d   s o   t h a t   f r a c t u r e   d o e s   n o t   o c c u r .   P e l l e t   X I I 1   f o r   w h i c h   t r u e   k i n e t i c   d a t a   h a d   b e e n   o b t a i n e d   w a s   m o u n t e d   a n d   p o l i s h e d   f o r   m e t a l l o g r a p h i c   e x a m i n a t i o n .   F i g u r e   l a   s h o w s   t h e   s p e c i m e n   a t   a   m a g n i f i c a t i o n   o f   c a .   4 X .   T h e   g r a y   o u t e r   r i m   i s   t h e   o x i d e ,   a n d   t h e   i n n e r   w h i t e   c i r c u l a r   a r e a   i s   t h e   s u r f a c e   o f   c a r b i d e .   T o   t h e   n a k e d   e y e ,   t h e   o u t e r   o x i d e   c o a t i n g   l o o k e d   w h i t e   a n d   c h a l k y ,   a n d   t h e   i n n e r   s u r f a c e   f r o m   w h i c h   t h e   o x i d e   c o a t i n g   w a s   p o l i s h e d   o f f ,   l o o k e d   b r i g h t   a n d   m e t a l l i c .   A   m o t t l e d   r i m   i s   c l e a r l y   v i s i b l e   a l o n g   t h e   o x i d e a l l o y   i n t e r f a c e   i n   t h e   f i g u r e   a n d   t h e   r e m a i n i n g   p h o t o m i c r o   g r a p h s   f o c u s   o n   p o r t i o n s   o f   t h i s   i n t e r f a c e .   I n   F i g .   l b ,   a t   a   m a g n i f i c a t i o n   o f   c a .   3 0 X ,   t h e   o x i d e   f i l l s   t h e   e n t i r e   u p p e r   h a l f   o f   t h e   p h o t o g r a p h ,   a n d   t h e   a l l o y   t h e   l o w e r .   T h e   o x i d e   i s   o b v i o u s l y   g r o w i n g   p r e f e r e n t i a l l y   a l o n g   g r a i n   b o u n d a r i e s   i n   t h e   c a r b i d e ,   a n d   e n v e l o p i n g   i n   d i v i d u a l   a l l o y   c r y s t a l l i t e s .   T h e   s t r u c t u r e   o f   t h e   b u l k   o x i d e   i s   v e r y   d i f f e r e n t   f r o m   t h a t   o n   Z r B 2   ( 8 ) ,   a l t h o u g h   a f t e r   c o o l i n g   b o t h   s h o w e d   o n l y   t h e   x r a y   l i n e s   f o r   m o n o c l i n i c   Z r O 2 .   O n   Z r B 2   t h e   o x i d e   w a s   s e e n   t o   g r o w   i n   a   c o l u m n a r   s t r u c t u r e ;   o n   Z r C ,   t h e   Z r O 2   a s s u m e d   a n   e q u i a x e d   g r a i n   s t r u c t u r e   v e r y   s i m i l a r   t o   t h a t   o f   t h e   o r i g i n a l   a l l o y .   F i g u r e s   l c   a n d   l d   s h o w   t h e   o x i d e   a l l o y   i n t e r f a c e   a t   a   s t i l l   h i g h e r   m a g n i f i c a t i o n ,   c a .   1 9 5 X ,   a n d   o n e   s e e s   e v e n   m o r e   c l e a r l y   t h e   p r e f e r e n t i a l   o x i d a t i o n   o f   g r a i n   b o u n d a r i e s ,   a n d   t h e   l a t e r a l   f i n g e r l i k e   g r o w t h   o f   o x i d e   f r o m   t h e   b o u n d a r i e s   i n t o   t h e   c r y s t a l l i t e   b u l k .   T h e   r e a c t i o n   z o n e   o f   i n t e r g r a n u l a r   a t t a c k   w a s   a p   p r o x i m a t e l y   0 . 0 1 4   + _   0 . 0 0 2   c m   i n   t h i c k n e s s .   T h e   o u t e r   o x i d e   w a s   a b o u t   t e n   t i m e s   t h i c k e r .   T h e   m e c h a n i s m   o f   o x i d e   g r o w t h   o n   Z r C   a t   h i g h   t e m p e r a t u r e s   a p p e a r s   t o   b e   r a p i d   a t t a c k   a t   g r a i n   b o u n d a r i e s ,   a n d   s l o w   o x i d a   t i o n   o f   t h e   a l l o y   f r o m   t h e   g r a i n   b o u n d a r y   s u r f a c e   i n   w a r d .   F i g u r e   l e ,   a t   a   m a g n i f i c a t i o n   o f   c a .   3 4 5 X ,   s h o w s   p o r t i o n s   o f   t h e   a l l o y   c o m p l e t e l y   e n v e l o p e d   b y   o x i d e .   F i n a l l y ,   F i g .   I f   i s   a   v i e w   o f   t h e   a l l o y   s u r f a c e   a t   a   m a g   n i f i c a t i o n   o f   c a .   3 4 5 X .   T h i s   s u r f a c e   h a d   b e e n   c o v e r e d   w i t h   a   d e n s e   o x i d e   p r i o r   t o   p o l i s h i n g ,   a n d   o n e   s e e s   h e r e   t h e   p e n e t r a t i o n   o f   t h e   o x i d e   i n t o   g r a i n   b o u n d a r i e s   o f   t h e   a l l o y .   B e t w e e n   1 1 2 6   ~   a n d   1 5 5 9 ~   t h e   g r a i n   b o u n d a r y   a t t a c k   r e s u l t s   i n   i n t e r g r a n u l a r   f r a c t u r e   o f   t h e   a l l o y .   A b o v e   1 5 8 0 ~   t h e r e   i s   a p p a r e n t l y   e n o u g h   p l a s t i c i t y   i n   e i t h e r   a l l o y ,   o x i d e ,   o r   b o t h ,   s o   t h a t   t h e   s a m p l e   r e m a i n s   i n t a c t   d u r i n g   o x i d a t i o n .   H a f n i u m   c a r b i d e . O x i d a t i o ~   k i n e t i c s . T h e   e x p e r i   m e n t a l   d a t a   f o r   H f C   a r e   s u m m a r i z e d   i n   T a b l e   I I .   S i n c e   t h e   a r c m e l t e d   s a m p l e s   w e r e   h i g h l y   i r r e g u l a r   i n   s h a p e ,   F i g .   1   a d .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   Z r C   ( p e l l e t   X I I 1 ) ;   T   1 9 6 ~   P o 2   2 5 . 9   T a r r .   M a g n i f i c a t i o n :   a   ( t a p   l e f t )   c a .   4 X ;   b   ( t a p   r i g h t )   c a .   3 0 X ;   c   ( b o t t o m   l e f t )   c a .   1 9 5 X ;   d   ( b o t t o m   r i g h t )   c a .   1 9 5 X .   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   1 0 3 2   J .   E l e c t r o c h e m .   S o c . :   E L E C T R O C H E M I C A L   S C I E N C E   O c t o b e r   1 9 6 7   F i g .   1   e f .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   Z r C   ( p e l l e t   X I I 1 ) ;   T   1 9 6 ~   P o 2   2 5 . 9   T o r r .   M a g n i f i c a t i o n :   e   ( l e f t )   a n d   f   ( r i g h t )   c a .   3 4 5 X .   F i g .   2 .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   h a f n i u m   c a r b i d e   p e l l e t   X V I 1 9 ;   t e m p e r a t u r e   1 6 0 0 ~ 1 7 6   P o 2   1 1 . 5   T o r r .   M a g n i f i c a t i o n   c a .   1 3 0 X .   t h e   g e o m e t r i c a l l y   c a l c u l a t e d   s u r f a c e   a r e a s   a r e   o n l y   a p p r o x i m a t e .   T h e   s u r f a c e   o x i d e   t e n d e d   t o   f l a k e   a n d   s p a l l   o n   r e m o v a l   o f   t h e   s a m p l e   f r o m   t h e   a p p a r a t u s ;   a n d ,   h e n c e ,   t h e   m e a s u r e d   w e i g h t   g a i n s   a r e   m i n i m u m   v a l u e s .   F r o m   t h e   m e a s u r e d   m i n i m u m   n e t   w e i g h t   g a i n s ,   a n d   t h e   w e i g h t s   o f   C O ( g )   a n d   C O s ( g )   i n   t h e   p r o d u c t   g a s   s t r e a m ,   t h e   m i n i m u m   r a t i o   o f   h a f n i u m   t o   c a r b o n   o x i d i z e d   c a n   b e   c a l c u l a t e d ,   a s   d e s c r i b e d   a b o v e   f o r   t h e   o x i d a t i o n   o f   Z r C .   T h e   e x p e r i m e n t a l   v a l u e s   o f   t h e   r a t i o   a r e   g i v e n   i n   T a b l e   I I   a s   1 4 . 1 ,   1 0 . 1 ,   a n d   1 5 . 9 ,   t o   b e   c o m   p a r e d   t o   a n   H f : C   r a t i o   i n   t h e   o r i g i n a l   a l l o y   o f   1 5 . 6 .   L i n e a r   o x i d a t i o n   w a s   o b s e r v e d   i n   e v e r y   c a s e .   F l o w   r a t e s   o f   1 1 9   c c / m i n ,   a p p r o x i m a t e l y   d o u b l e   t h o s e   u s e d   f o r   t h e   Z r C   e x p e r i m e n t s ,   w e r e   e m p l o y e d   i n   o r d e r   t o   m a i n t a i n   a n   a d e q u a t e   s u p p l y   o f   o x y g e n .   A t   1 2 8 0 ~   i n   p u r e   o x y g e n   a t   1   a t m ,   a   h a f n i u m   c a r   b i d e   s p e c i m e n   d i s i n t e g r a t e d   i n t o   s e v e r a l   p i e c e s   w i t h i n   3   m i n   i n   a   m a n n e r   v e r y   s i m i l a r   t o   t h a t   d e s c r i b e d   a b o v e   f o r   z i r c o n i u m   c a r b i d e .   M e t a l l o g r a p h i c   e x a m i n a t i o n   o f   o x i d e   f i l m s . ~ P e l l e t   X V I 1 9   w a s   i m b e d d e d   i n   p l a s t i c   a n d   p o l i s h e d   f o r   m i   c r o s c o p i c   o b s e r v a t i o n .   I n   t h e   p h o t o m i c r o g r a p h s   s h o w n   i n   F i g .   2 ,   t h e   o x i d e   i s   s e e n   t o   c o n t a i n   c o n s i d e r a b l e   p o   r o s i t y .   A s   i n   t h e   c a s e   o f   z i r c o n i u m   c a r b i d e ,   o x i d a t i o n   a p p e a r s   t o   b e   p r e f e r e n t i a l   a l o n g   g r a i n   b o u n d a r i e s .   T h e   p h o t o m i c r o g r a p h   i n   F i g .   3   i s   f o r   a   h a f n i u m   c a r   b i d e   p e l l e t   t h a t   h a d   b e e n   e x p o s e d   t o   a   N 2 H e   m i x t u r e   f o r   a n   h o u r   a n d   a   q u a r t e r   a t   a   t e m p e r a t u r e   o f   I 9 6 0 ~   T h e   N u   a p p a r e n t l y   c o n t a i n e d   a   s m a l l   q u a n t i t y   o f   o x y   g e n ,   a n d   w h i t e   H f O 2 ( c )   w a s   i d e n t i f i e d   o n   t h e   s u r f a c e   o f   t h e   s a m p l e   b y   x r a y   d i f f r a c t i o n .   T h e   n e t   w e i g h t   c h a n g e   o f   t h e   s a m p l e   w a s   0 . 0 0 2 4   g / c m   2 ,   a n   o r d e r   o f   m a g n i t u d e   l e s s   t h a n   t h a t   o b s e r v e d   a t   a n   o x y g e n   p a r   t i a l   p r e s s u r e   o f   1 1 . 5   T o r r   a t   a b o u t   t h e   s a m e   t e m p e r a   t u r e .   T h e   o x i d a t i o n   r a t e   w a s   s e e n   t o   d e c r e a s e   s l i g h t l y   w i t h   t i m e ,   a n d ,   a l t h o u g h   g r a i n   b o u n d a r y   o x i d a t i o n   i s   T a b l e   I I .   S u m m a r y   o f   F i g .   3 .   P h o t o m i c r o g r a p h   a f   o x i d i z e d   h a f n i u m   c a r b i d e   p e l l e t   X V I 2 2 ;   t e m p e r a t u r e   1 9 6 0 ~   l o w   o x y g e n   p r e s s u r e .   M a g n i f i c a t i o n   c a .   2 3 0 X .   a p p a r e n t   i n   F i g .   3 ,   t h e   o x i d e   s e e m s   l e s s   p o r o u s   t h a n   t h a t   s h o w n   i n   F i g .   2 ,   f o r   o x i d a t i o n   a t   a   h i g h e r   p r e s   s u r e .   D i s c u s s i o n   A   r e c e n t   r e p o r t   ( 9 )   s u g g e s t s   t h a t   g r a i n   b o u n d a r y   o x i d a t i o n   i s   c h a r a c t e r i s t i c   O f   i r o n   c o n t a m i n a t e d   Z I C .   F o r   m a t e r i a l   w i t h   a n   i m p u r i t y   l e v e l   a b o u t   1   v / o   F e s C   o r   a b o u t   1 . 1   w / o ,   t h e   F e 3 C   i s   f o u n d   t o   b e   s e g r e g a t e d   a t   t h e   g r a i n   b o u n d a r i e s ,   a n d   t o   b e   o x i d i z e d   a t   a   m u c h   m o r e   r a p i d   r a t e   t h a n   t h e   Z r C   m a t r i x .   A   c h e m i c a l   a n a l y s i s   o f   t h e   z o n e   m e l t e d   m a t e r i a l   u s e d   i n   t h e   p r e s e n t   s t u d y   s h o w e d   a n   i r o n   c o n t a m i n a   t i o n   l e v e l   o f   0 . 0 7 %   b y   w e i g h t .   I n   o r d e r   t o   t r y   t o   a s s e s s   t h e   p o s s i b l e   i n f l u e n c e   o f   i r o n   o n   t h e   p r e s e n t   r e s u l t s ,   t h e   o x i d i z e d   a n d   p o l i s h e d   Z r C   s a m p l e   s h o w n   i n   F i g .   1   w a s   e x a m i n e d   f o r   i r o n   w i t h   t h e   e l e c t r o n   p r o b e .   1   A n   F e K ~ I   s c a n n i n g   i m a g e   w a s   t a k e n   o f   t h e   s a m p l e ,   a n d   i r o n   w a s   f o u n d   t o   b e   c o n c e n t r a t e d   i n   t h e   s m a l l   c i r c u l a r   i n c l u s i o n s   v i s i b l e   i n   F i g .   1 ,   a n d   i n h o m o g e n e   o u s l y   i n   t h e   c r a c k s .   P o i n t   c o u n t   a n a l y s e s   t a k e   i n   t h e   i n c l u s i o n s ,   t h e   c r a c k s ,   t h e   g r a i n   b o u n d a r i e s ,   a n d   t h e   m a t r i x   s h o w e d   t h e   p r e s e n c e   o f   i r o n   p a r t i c l e s   o f   l e s s   t h a n   1 ~   i n   s i z e   d i s t r i b u t e d   a t   r a n d o m   i n   t h e   i n c l u s i o n s   a n d   c r a c k s ,   b u t   f a i l e d   t o   r e v e a l   t h e   p r e s e n c e   o f   a n y   i r o n   a t   a l l   i n   e i t h e r   t h e   g r a i n   b o u n d a r i e s   o r   t h e   m a t r i x .   T h e   p r e f e r e n t i a l   g r a i n   b o u n d a r y   o x i d a t i o n   o b s e r v e d   i n   z i r c o n i u m   c a r b i d e ,   t h e r e f o r e ,   s e e m s   t o   b e   c h a r a c t e r   i s t i c   o f   t h e   p u r e   m a t e r i a l .   A c k n o w l e d g m e n t   I t   i s   a   p l e a s u r e   t o   t h a n k   J o h n   E n g e l k e   o f   A r t h u r   D .   L i t t l e ,   I n c . ,   f o r   h i s   a s s i s t a n c e   i n   o b t a i n i n g   a n d   i n t e r   p r e t i n g   t h e   e l e c t r o n   p r o b e   d a t a ,   a n d   t o   a c k n o w l e d g e   t h e   i n v a l u a b l e   a s s i s t a n c e   o f   R i c h a r d   F .   Q u i g l e y   a n d   W a l t e r   C h r i s t e n s e n   w i t h   t h e   e x p e r i m e n t a l   w o r k .   M a n u s c r i p t   r e c e i v e d   J a n .   2 3 ,   1 9 6 7 ;   r e v i s e d   m a n u   s c r i p t   r e c e i v e d   M a y   3 1 ,   1 9 6 7 .   A n y   d i s c u s s i o n   o f   t h i s   p a p e r   w i l l   a p p e a r   i n   a   D i s c u s s i o n   S e c t i o n   t o   b e   p u b l i s h e d   i n   t h e   J u n e   1 9 6 8   J O U R N A L .   1 T h e   e l e c t r o n   p r o b e   a n a l y s i s   w a s   k i n d l y   s u p p l i e d   b y   A c t o n   L a b o r a t o r i e s ,   5 3 1   M a i n   S t r e e t ,   A c t o n ,   M a s s a c h u s e t t s .   e x p e r i m e n t a l   d a t a   o n   H f C   I n i t i a l   P r e s s u r e ,   T i m e ,   E x p t .   w e i g h t ,   g   A r e a ,   c m   ~   T e m p ,   \" K   T o r r   m i n   M i n i m u m   M i n i m u m   C   H f   n e t   w e i g h t   C O   C O s   c o n s u m e d ,   c o n s u m e d ,   H f / C   c h a n g e ,   g   f o r m e d ,   g   f o r m e d ,   g   g   g   r a i n   X V I I 4   0 . 5 3 7 4   1 , 0 7 6   1 7 9 0   1 1 . 5   8 5   X V I 3 8   0 . 6 5 4 4   1 . 2 4 1   1 8 9 0   1 1 . 5   8 9   X V I 3 6   0 . 7 5 1 1   1 . 0 3 2   2 0 0 5   1 1 . 5   4 3   X V I 1 9   0 . 6 0 5 7   1 . 2 5 6   2 0 0 0   7 . 5   1 2 0   1 6 0 0   ( c o n t i n u o u s   d r o p )   0 . 0 2 1   0 . 0 1 6 0   0 . 0 2 5 4   0 . 0 1 3 8   0 . 1 9 5   1 4 . 1   0 . 0 1 6   0 . 0 2 3 2   0 . 0 3 5 5   0 . 0 1 9 6   0 . 1 9 9   1 0 . 1   0 . 0 2 2     0 . 0 1 9 1   0 , 0 3 8   0 . 0 3 4 0   0 . 0 2 2 5   0 . 0 2 0 7   0 , 3 2 8   1 5 . 9   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   V o l .   1 1 4 ,   N o .   1 0   H I G H T E M P E R A T U R E   O X I D A T I O N   1 0 3 3   R E F E R E N C E S   1 .   A .   K .   K u r i a k o s e   a n d   J .   L .   M a r g r a v e ,   T h i s   J o u r n a l ,   1 1 1 ,   8 2 7   ( 1 9 6 ~ t ) .   2 .   R .   W .   B a r t l e t t ,   M .   E .   W a d s w o r t h ,   a n d   I .   B .   C u t   l e r ,   T r a n s .   A I M E ,   2 2 7 ,   4 6 7   ( 1 9 6 3 ) .   3 .   W .   W a t t ,   G .   H .   C o c k e t t ,   a n d   A .   R .   H a l l ,   M e t a u x ,   2 8 ,   2 2 2   ( 1 9 6 3 ) .   4 .   E .   F .   W e s t r u m ,   J r . ,   a n d   G .   F e i c k ,   J .   C h e m .   E n g .   D a t a ,   8 ,   1 7 6   ( 1 9 6 3 ) .   5 .   L .   A .   M c C l a i n e ,   T h e r m o d y n a m i c   a n d   K i n e t i c   S t u d i e s   f o r   a   R e f r a c t o r y   M a t e r i a l s   P r o g r a m ,   4 t h   S e m i   a n n u a l   P r o g r e s s   R e p o r t   ( A u g u s t ,   1 9 6 3 ) ,   C o n t r a c t   N o .   A F   3 3   ( 6 1 6 ) 7 4 7 2 .   6 .   J .   B .   B e r k o w i t z N I a t t u c k ,   T h i s   J o u r n a l ,   I l l ,   9 0 8   ( 1 9 6 4 ) .   7 .   J .   B .   B e r k o w i t z M a t t u c k   a n d   R .   R .   D i l s ,   i b i d . ,   1 1 2 ,   5 8 3   ( 1 9 6 5 ) .   8 .   J .   B .   B e r k o w i t z M a t t u c k ,   i b i d . ,   1 1 3 ,   9 0 8   ( 1 9 6 6 ) .   9 .   K .   R .   J a n o w s k i ,   R .   D .   C a r n a h a n ,   a n d   R .   C .   R o s s i ,   \" S t a t i c   a n d   D y n a m i c   O x i d a t i o n   o f   Z r C , \"   T D R   6 6 9   ( 6 2 5 0 1 0 ) 3 ,   A e r o s p a c e   C o . ,   E l   S e g u n d o ,   C a l i f . ,   J a n u a r y ,   1 9 6 6 .   F u n d a m e n t a l   L i m i t a t i o n s   o n   t h e   L o w T e m p e r a t u r e   O p e r a t i o n   o f   E l e c t r o l y t i c   D e v i c e s   C .   A .   A n g e l l   D e p a r t m e n t   o f   C h e m i s t r y ,   P u r ~ u e   U n i v e r s i t y ,   L a f a y e t t e ,   I n d i a n a   A B S T R A C T   A   n e w   a p p r o a c h   t o   t h e   u n d e r s t a n d i n g   o f   t r a n s p o r t   b e h a v i o r   i n   c o n c e n   t r a t e d   e l e c t r o l y t e   s o l u t i o n s   r e q u i r e s   t h e   r e c o g n i t i o n   a t   l o w   t e m p e r a t u r e s   o f   a   l i q u i d   s t a t e   l i m i t i n g   t e m p e r a t u r e   w h i c h   i s   a   t h e r m o d y n a m i c   c o n s t a n t   o f   a n y   s o l v e n t e l e c t r o l y t e   s o l u t i o n .   T h e   v a l u e   o f   t h i s   c o n s t a n t ,   w h i c h   d e t e r m i n e s   s e r v i c e   l i m i t s ,   i s   a   f u n c t i o n   o f   e l e c t r o l y t e   c o n c e n t r a t i o n   e x c e p t   a t   l o w   c o n   c e n t r a t i o n s   w h e r e   s o l v e n t   s t r u c t u r e   m a y   d o m i n a t e .   I n   t h i s   p a p e r   t h e   r e s u l t s   o f   w o r k e r s   i n v e s t i g a t i n g   e l e c t r o l y t e   s o l u t i o n s   f o r   l o w t e m p e r a t u r e   b a t t e r y   a p p l i c a t i o n s   a r e   u s e d   t o   e m p h a s i z e   t h e   u s e f u l n e s s   o f   t h e   c o n c e p t s   u n d e r l y i n g   t h i s   a p p r o a c h   t o   l o w t e m p e r a t u r e   e l e c t r o l y t e   p r o b l e m s .   T h i s   a r t i c l e   c o n s i d e r s   b r i e f l y   s o m e   i m p l i c a t i o n s   o f   r e c e n t   d e v e l o p m e n t s   i n   e l e c t r o l y t e   s o l u t i o n   t r a n s p o r t   t h e o r y   t o   t h e   p r a c t i c a l   p r o b l e m s   e n c o u n t e r e d   i n   t h e   u s e   o f   e l e c t r o l y t e   s o l u t i o n s   a t   l o w   t e m p e r a t u r e s .   T e m p e r a t u r e   a n d   c o n c e n t r a t i o n   d e p e n d e n c e   o f   s o t u   t i o n   t r a n s p o r t   p r o p e r t i e s . I n   r e c e n t   p a p e r s   ( 1 3 )   i t   h a s   b e e n   s h o w n   t h a t   t h e   t e m p e r a t u r e   d e p e n d e n c e   a n d   c o m p o s i t i o n   d e p e n d e n c e   o f   e l e c t r i c a l   ( e q u i v a l e n t )   c o n d u c t a n c e ,   A ,   a n d   v i s c o u s   f l o w ,   ~ ] ,   p r o c e s s e s   i n   v a r i   o u s   c o n c e n t r a t e d   a q u e o u s   e l e c t r o l y t e   s o l u t i o n s   a t   l o w   c o r r e s p o n d i n g   t e m p e r a t u r e s   m a y   b e   d e s c r i b e d   b y   e q u a t i o n s   o f   t h e   f o l l o w i n g   f o r m   c o n s t a n t   c o m p o s i t i o n   k   A ( I ~ )   ,   1 / ~ ] ( N I   =   A   e x p   .   T     T o   c o n s t a n t   t e m p e r a t u r e   k / Q   A ( T )   ,   1 / ~ ] ( T )   . ~   A   e x p   N o     N   [ 1 ]   [ 2 ]   w h e r e   A ,   k ,   a n d   T o   i n   E q .   [ 1 ]   a n d   A ,   k / Q ,   a n d   N o   i n   E q .   [ 2 ]   a r e   c o n s t a n t s ;   T   a n d   N   a r e ,   r e s p e c t i v e l y ,   t h e   a b s o l u t e   t e m p e r a t u r e   a n d   t h e   e q u i v a l e n t   c o n c e n t r a t i o n .   A   p l a u s i b l e   t h e o r e t i c a l   i n t e r p r e t a t i o n   d i s c u s s e d   i n   d e t a i l   e l s e w h e r e   ( 1 ,   4 ,   5 )   m a y   b e   g i v e n   t h e s e   e q u a t i o n s   i n   t e r m s   o f   t h e   c o n t r o l l i n g   i n f l u e n c e   o n   t h e   t r a n s p o r t   p r o c e s s e s ,   o f   t h e   l i q u i d   c o n f i g u r a t i o n a l   e n t r o p y   c o n   t e n t   ( 6 ) .   T h e   c o n s t a n t s   T o   a n d   N o   i n   t h i s   c a s e   r e p r e   s e n t   t h e   t e m p e r a t u r e   a t   f i x e d   c o m p o s i t i o n ,   a n d   t h e   e q u i v a l e n t   c o n c e n t r a t i o n   a t   c o n s t a n t   t e m p e r a t u r e ,   r e   s p e c t i v e l y ,   a t   w h i c h   t h e   c o n f i g u r a t i o n a l   e n t r o p y   v a n i s h e s .   S i g n i f i c a n c e   o . f   l o w t e m p e r a t u r e   e l e c t r o l y t e   s o t u   t i o n   b e h a v i o r . A n   i m p o r t a n t   r e s u l t   o f   t h i s   t r e a t m e n t   w h i c h   w e   w i s h   t o   e m p h a s i z e   i n   t h i s   p a p e r   i s   t h e   r e a l i z a t i o n   t h a t ,   p r o v i d e d   c r y s t a l l i z a t i o n   h a s   n o t   a l   r e a d y   o c c u r r e d ,   a n y   e l e c t r o l y t e   s o l u t i o n   w i l l   l o s e   i t s   l i q u i d   c h a r a c t e r   a n d   b e c o m e   a   g l a s s   a t   a   ~ e m p e r a t u r e   ( ~ T o )   w h i c h   i n   g e n e r a l   w i l l   b e   w e l l   a b o v e   1 0 0 ~   W h e r e   t h e   e l e c t r o l y t e   p r o p e r t y   o f   i n t e r e s t   d e p e n d s   o n ,   o r   i s   r e l a t e d   t o ,   t h e   f l u i d i t y   o f   t h e   s u b s t a n c e ,   t h i s   t e m   p e r a t u r e   t h e r e f o r e   p l a c e s   a n   a b s o l u t e   l o w e r   l i m i t   t o   t h e   s e r v i c e a b i l i t y   o f   t h e   m a t e r i a l .   1   F u r t h e r m o r e ,   i t   i s   f o u n d   t h a t ,   f o r   m a n y   e l e c t r o l y t e   m i x t u r e s   a n d   s o l u t i o n s   a t   t e m p e r a t u r e s   n o t   t o o   f a r   a b o v e   T o ,   T o   i t s e l f   i s   t h e   o n l y   i m p o r t a n t   v a r i a b l e   i n   t h e   t r a n s p o r t   e q u a t i o n ,   s o   t h a t   t h e   l o w t e m p e r a t u r e   t r a n s p o r t   p r o p e r t i e s   o f   s u c h   l i q u i d s   a r e   t o   a   l a r g e   e x   t e n t   k n o w n   o n c e   T o   i s   k n o w n   ( 3 ) .   T h u s ,   t h e   s c a l i n g   f a c t o r   t o   b e   u s e d   i n   c o m p a r i n g   a   g i v e n   p r o p e r t y   a m o n g s t   d i f f e r e n t   s o l u t i o n s   i s ,   f r o m   E q .   [ 1 ] ,   ( T   T o ) .   T o   i l l u s t r a t e   t h e   u s e f u l n e s s   o f   t h i s   p o i n t ,   w e   t a k e   a n   e x a m p l e   f r o m   t h e   w o r k   o f   G a r r e t t   e t   a l .   ( 7 )   w h o   w e r e   i n v e s t i g a t i n g   v a r i o u s   s t r o n g   ( p r e s u m e d   f u l l y   d i s s o   c i a t e d )   e l e c t r o l y t e   s o l u t i o n s   f o r   s u i t a b i l i t y   i n   l o w   t e m p e r a t u r e   b a t t e r y   a p p l i c a t i o n s .   T h e i r   ( u n i n t e r   p r e t e d )   d a t a   o n   t h e   v i s c o s i t y   o f   s o m e   s o l u t i o n s ,   r e l a   t i v e   t o   t h e   v i s c o s i t y   o f   w a t e r   a t   2 5 ~   a r e   r e p r o d u c e d   i n   F i g .   l ( i ) .   T h e   s t r i k i n g   f e a t u r e   o f   t h e   d a t a   i s ,   o f   c o u r s e ,   t h e   v e r y   r a p i d   r i s e   i n   s o l u t i o n   v i s c o s i t y   a t   t h e   l o w e r   t e m p e r a t u r e s .   I n   F i g .   1   ( i i )   w e   s h o w   h o w   t h e   u s e   o f   a   T o   v a l u e   a p p r o p r i a t e   t o   e a c h   s o l u t i o n   i n   t h e   s c a l   i n g   f a c t o r   ( T   T o )   r e d u c e s   t h e   d a t a   a p p r o x i m a t e l y   t o   a   s i n g l e   c u r v e .   T h e   r e m a i n i n g   m i n o r   d i f f e r e n c e s   s e e m   t o   b e   d u e   m a i n l y   t o   v a r i a t i o n s   i n   t h e   v a l u e   o f   t h e   p r e e x p o n e n t i a l   t e r m   A ,   a s   s e e n   i n   t h e   f o l l o w i n g .   A c c o r d i n g   t o   E q .   [ 1 ] ,   t h e   v a l u e s   o f   T o   w h i c h   r e d u c e   t h e   s o l u t i o n   v i s c o s i t i e s   a s   i n   F i g .   l ( i i ) ,   s h o u l d   y i e l d   a   l i n e a r   p l o t   f o r   t h e   r e l a t i v e   v i s c o s i t i e s   w h e n   l o g   ~ / ~ ] o ,   ( o r   l o g   ~ o / ~   t o   m a i n t a i n   E q .   [ 1 ]   s i g n s )   i s   p l o t t e d   a g a i n s t   1 / ( T T o ) .   T h e   a p p r o p r i a t e   s e m i l o g a r i t h m i c   p l o t s   a r e   s h o w n   i n   F i g .   2 .   T h e   v a r i o u s   p l o t s   a r e   n o w   s e e n   t o   b e   d i f f e r e n t i a t e d   b y   s m a l l   c h a n g e s   i n   t h e   v a l u e s   o f   t h e   p a r a m e t e r   A ,   t h e   p l o t s   b e i n g   l i n e a r   w i t h   e s   s e n t i a l l y   e q u a l   v a l u e s   o f   t h e   s l o p e   k .   I t   m u s t   b e   s a i d ,   h o w e v e r ,   t h a t   t h e   a v a i l a b l e   d a t a   a r e   n o t   s u f f i c i e n t l y   T h e   l i m i t i n g   t e m p e r a t u r e   u s u a l l y   i m p o s e d   b y   t h e   c r y s t a l l i z a   t i o n   t e m p e r a t u r e   i s   n o t   a n   a b s o l u t e   l i m i t   i n s o f a r   a s   s u i t a b l e   a d d i   t i v e s   c a n   u s u a l l y   m a k e   c r y s t a l   n u c l e a t i o n   a   v e r y   i m p r o b a b l e   p r o c e s s   e v e n   w h e n   t h e r m o d y n a m i c a l l y   f a v o r e d .   \\x0c']"
},{
  "_id": 92,
  "PDF": "High-Temperature Oxidation of ZrB2–MoSi2–AlN Composite Ceramics.pdf",
  "Text": "['DOI 10.1007/s11106-019-00052-5  Powder Metallurgy and Metal Ceramics, Vol. 58, Nos. 1-2, May, 2019 (Russian Original Vol. 58, Nos. 1-2, Jan.-Feb., 2019)   HIGH-TEMPERATURE OXIDATION   OF ZrB2-MoSi2-AlN COMPOSITE CERAMICS   O.N. Grigoriev,1 A.D. Panasyuk,1 I.A. Podchernyaeva,1,2   I.P. Neshpor,1 and D.V. Yurechko1   UDC 621.762;620.93   The oxidation mechanism of ZrB2-MoSi2-AlN composite ceramics has been studied in the range  1550-1680\\uf0b0C. The oxidation  stages and  scale  structure and phase composition have been  established. Two main scale layers have been found.  The outer layer forms at \\uf07e1550\\uf0b0C. It consists  of SiO2 (\\uf062-cristobalite) and an Al2O3-SiO2 solid solution containing MoO3. At 1680\\uf0b0C, a layer of   crystalline zirconium dioxide grains forms at the boundary with the initial ceramic surface. This   layer is an effective protective barrier for oxygen diffusion into the sample.   Keywords:   zirconium   diboride, molybdenum   disilicide,   aluminum   nitride,   ceramics,   high temperature oxidation.   INTRODUCTION   To develop corrosion-resistant ultrahigh-temperature ceramics (UHTC), silicon-containing additions, such  as SiC, ZrSi2, MoSi2, etc., are introduced into the ZrB2 matrix. These additions lead to a silicon dioxide film on the  scale that is thermodynamically stable up to \\uf07e1600°C [1-5]. The composition of the film, which is a barrier to the   diffusion of oxygen into the sample, can be controlled by the introduction of complex additives.    As we showed previously [7], the desired effect is reached with the use of a complex AlN-SiC addition,  producing a thin film of Al2O3-SiO2 mullite solid solutions that are thermodynamically stable up to temperatures   above 1700°C.   The oxidation mechanism of ZrB2-MoSi2 UHTC in a wide MoSi2 range (14-44 wt.%) was examined; it   was shown that this system was promising for the development of alloys with an increased content of a specific  addition [2]. The desired effect is reached through the structurization of the surface film in which the outer SiO2  layer forms on the Mo5Si3 particles, preventing the oxidation of MoSi2 and Mo5Si3.   The surface film is reinforced with ZrO2 grains adhesively bonded to the oxide layer. In this regard, it is of  interest to examine the mechanism of high-temperature oxidation of ZrB2-based UHTC with a complex MoSi2-AlN   addition.  The objective of this paper is to examine the oxidation mechanism of ZrB2-MoSi2 ultrahigh-temperature   ceramics with an AlN addition in the temperature range to 1550-1700°C.   1Frantsevich Institute for Problems of Materials Science, National Academy of Sciences of Ukraine, Kyiv,   Ukraine.     2To whom correspondence should be addressed; e-mail: iripodcher@gmail.com.   Translated from Poroshkova Metallurgiya, Vol. 58, Nos. 1-2 (525), pp. 124-129, 2019. Original article   submitted June 26, 2018.       1068-1302/19/0102-0099 \\uf0d32019 Springer Science+Business Media, LLC                            99\\xa0  \\x0c', 'EXPERIMENTAL PROCEDURE   Compact samples of composition ZrB2-20MoSi2-10AlN (here and further in wt.%) were produced by hot   pressing [6] from ultrafine powders (1-5 µm average size) at 1850°C and 30 MPa for ~15 min followed by holding   for 35 min at 1350°C. The samples were oxidized in air at 1550°C for 3 h employing a Nabertherm 4HT unit.   The structure and elemental composition of the oxidized samples were examined on ZrB2-20MoSi2-10AlN   cross-sections. Their microstructure and quantitative elemental composition were studied using a REM-106   microscope equipped with an energy-dispersive chemical analysis system.   RESULTS AND DISCUSSION   We previously [2] examined the high-temperature oxidation mechanism of ZrB2-MoSi2 UHTC (at 14- 44 wt.% MoSi2) in air at a heating rate of 15 min to ~1680°C. We established oxidation stages beginning from 950-  1100°C up to 1680°C. At the first stage (at ~950°C), zirconium boride oxidizes by the following reaction:    2ZrB2 + 5O2 = 2ZrO2 + 2B2O5 \\uf0ad.   (1)   When temperature increases to \\uf03e1100°C, MoSi2 starts oxidizing to form thermodynamically stable Mo5Si3  (to temperatures above 1680°C) and an amorphous SiO2 film over the silicide phases:   5MoSi2 + 7O2 = Mo5Si3 + 7SiO2.   (2)   Reaction (2) proceeds with free silicon (\\uf07e6.2%) present in MoSi2 in the homogeneity range. This gives rise  to Mo5Si3, which is thermodynamically stable to \\uf07e1700\\uf0b0C, and SiO2. The SiO2 oxide film stabilizes at \\uf07e1700\\uf0b0C.  Silicon oxide forms an amorphous, partially discrete film that crystallizes into \\uf062-cristobalite (reaction (2)) at  \\uf07e1550\\uf0b0C. Boron oxide produced by reaction (1) dissolves in this oxide film to form borosilicate glass that may  contain individual ZrO2 grains [2]. The SiO2-based oxide film is weakly bonded to the scale (basic oxide layer) and   does not provide adequate protection against oxygen diffusion into the sample.   To increase oxidation resistance (to \\uf0b31800\\uf0b0C), we proposed the introduction of 10% AlN into the ZrB2- 20% MoSi2 composite. The cross-sectional microstructure of the ZrB2-20MoSi2-10AlN ceramic sample oxidized  at 1500\\uf0b0C for 1 h is shown in Fig. 1. Four layers were identified: an upper loose layer (1), outer and internal scale   layers (2, 3), and a base layer (4). The first three zones represent a scale oxide layer. The upper loose layer consists  of xAl2O3-ySiO2 solid solution (lower mullite) grains containing MoO3 as a solid solution (Fig. 2, point 2) and   individual zirconium oxide grains.  The outer scale layer (Fig. 2) has a heterophase structure and consists of the xAl2O3-ySiO2 solid solution  phase (Fig. 2, point 4) containing MoOx (Fig. 2, point 5); individual zirconium oxide grains are located along the   grain boundaries of the mullite solid solution (Fig. 2, point 3).  Therefore, the outer scale layer consists of complex oxide phases, such as individual Al2O3-SiO2 mullite  grains containing MoO3 as a solid solution (Fig. 2, point 5) and individual ZrO2 grains (Fig. 2, point 3). There is   strong adhesion between these phases and virtually zero porosity.   Fig. 1. Cross-sectional structure of the oxidized ZrB2-20MoSi2-10AlN composite   100       \\x0c', 'Fig. 2. Cross-sectional microstructure of upper loose layer 1 and inner scale layer 2 of the ceramic   sample and elemental composition (wt.%) of structural components   Fig. 3. Cross-sectional microstructure of the inner scale layer in the ceramic sample and elemental   composition of its structural components   The inner scale layer consists of zirconium oxide grains (Fig. 3, point 1). The Al-Si-O oxide phase is  located at the zirconium oxide grain boundaries and is an Al2O3-SiO2 solid solution (Fig. 3, point 2) containing  MoO3 and individual ZrO2 grains. It should be noted that there is adhesion bond between the phases and the oxide   grains might form in the presence of the liquid phase. The oxide layer is protective since the diffusion of silicon and  oxygen into the ceramic layer is virtually zero. The oxidation of MoSi2 proceeds by reaction (2): the resultant  silicon oxide protects Mo5Si3 and MoSi2 against oxidation.  Molybdenum oxides in mullite (Fig. 3, point 2) and, perhaps, silicide Mo5Si3 are revealed in the inner scale   layer. When oxidized, aluminum nitride promotes  the formation of mullite  increasing  the high-temperature  oxidation resistance of the ZrB2-MoSi2 composite up to \\uf07e1700\\uf0b0C. At the same time, mullites do not always   crystallize as needles and sometimes form isometric grains.   101                          \\x0c', 'Fig. 4. Cross-sectional microstructure of the starting ceramic sample in the ZrB2-20MoSi2-10AlN   system    The mullite lattice consists of chains formed by octahedrons of oxygen atoms around a part of aluminum   atoms connected by edges. The other part of aluminum and silicon atoms forms centers of oxygen atom  tetrahedrons. This also happens in the dissolution of other metal oxides (Cr2O3, Fe2O3, TiO2, etc.) forming solid  solutions. In our case, when a solid solution of molybdenum oxides in mullite forms, the Al2O3-SiO2 phase is close   to oval (Fig. 4).  The inner scale layer (Fig. 1, layer 3) borders the starting layer of the ZrB2-20MoSi2-10AlN composite   (Fig. 1, layer 4) whose structure is heterophase: the main phases are zirconium diboride, molybdenum disilicide  MoSi2, and a solid solution in the Al-N-Si-Mo3Si system. The solid solution grains are located between ZrB2 and  MoSi2 and form complex Al2O3-SiO2-MoO3 oxide in the oxidation process. Its shape is close to oval (Fig. 4).  To increase the high-temperature oxidation resistance of the composite ceramics to 1800-1850\\uf0b0C, a great  number of AlN and MoSi2 compounds needs to be introduced to activate sintering and produce high-temperature  oxide solid solutions in the oxidation process, which remain stable up to \\uf07e1900\\uf0b0C.   CONCLUSIONS   The oxidation mechanism of composite ceramics in the ZrB2-20MoSi2-10AlN (wt.%) system in the range  1550-1680\\uf0b0C has been studied.   At the first oxidation stage (at 1350\\uf0b0C), silicon present in MoSi2 in the homogeneity range (\\uf07e6.9 wt.%)  produces amorphous SiO2. This gives rise to thermodynamically stable Mo5Si3.   At \\uf07e1550\\uf0b0C, the oxide layer stabilizes through the formation of SiO2 \\uf062-cristobalite and an Al2O3-SiO2  solid solution containing MoO3 and individual ZrO2 grains. When oxidation temperature increases to \\uf07e1680\\uf0b0C, a   layer of zirconium oxide crystalline grains forms at the boundary with the starting ceramic surface and is an   effective protection barrier to the diffusion of oxygen into the sample.  Aluminum nitride performs the following functions in the ZrB2-MoSi2-AlN ceramic composites. When  oxidized, the Al-N-Si solutions form an Al2O3-SiO2 solid solution that activates the sintering process, determines   adhesion  interaction with molybdenum and zirconium oxides, and  resistance of the ZrB2-MoSi2-AlN composite to \\uf03e1700\\uf0b0C.   increases   the high-temperature oxidation   REFERENCES    O.N. Grigoriev, B.A. Galanov, V.A. Lavrenko, A.D. Panasyuk, and K.G. Nickel, “Oxidation of ZrB2-SiC- ZrSi2 ceramics in oxygen,” J. Ceram. Soc., 30, No. 11, 2397-2405 (2010).   V.A. Lavrenko, A.D. Panasyuk, O.M. Grigoriev, O.V. Koroteev, and V.A. Kotenko, “High-temperature (up  to 1600\\uf0b0C) oxidation of ZrB2-MoSi2 ceramics in air,” Powder Metall. Met. Ceram., 51, No. 1-2, 102-107   (2012).   1.   2.   102     \\x0c', '3.   4.   5.   6.   7.   H. Ping, X.-X. Zhang, J.-C. Han, X.-G. Luo, and S.-Y. Du, “Effect of various additives on the oxidation  behavior of ZrB2-based ultra-high-temperature ceramics at 1800\\uf0b0C,” J. Amer. Ceram. Soc., 93, No. 2, 345-  349 (2010).   S. Diletta, B. Mylene, and J. Alida, “Long-term oxidation, behavior and mechanical strength degradation of  a pressurelessly sintered ZrB2-MoSi2 ceramic,” Acta Mater., 53, 1297-1302 (2005).  T.Z. Vuntian, S. Marins, D.C. Samuel, and P.B. Davril, “Thermal oxidation kinetic is of MoSi2-based   powder,” J. Amer. Ceram. Soc., 82, 2785-2790 (1999).   O.N. Grigoriev, B.A. Galanov, A.V. Koroteev, L.M. Melakh, T.V. Mosina, and N.D. Bega,“Structurization   and mechanical properties of zirconium diboride in the presence of sintering activators,” in: Electron   Microscopy and Strength of Materials (Collected Papers) [in Russian], Inst. Probl. Materialoved. NAN   Ukrainy, No. 21, 111-129 (2015).   O.N. Grigoriev, A.D. Panasyuk, I.A. Podchernyaeva, I.P. Neshpor, and D.V. Yurechko, “Mechanism of  high-temperature oxidation of ZrB2-based composite ceramics in the ZrB2-SiC-AlN system,” Powder   Metall. Met. Ceram., 57, No. 1-2, 71-74 (2018).   103       \\x0c']"
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  "PDF": "High-Temperature Oxidation Zirconium and Hafnium Carbides.pdf",
  "Text": "['Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   H i g h T e m p e r a t u r e   O x i d a t i o n   I V .   Z i r c o n i u m   a n d   H a f n i u m   C a r b i d e s   J o a n   B .   B e r k o w i t z M a t t u c k   A r t h u r   D .   L i t t l e ,   I n c . ,   C a m b r i d g e ,   M a s s a c h u s e t t s   A B S T R A C T   T h e   o x i d a t i o n   o f   Z r C   w a s   s t u d i e d   a t   t e m p e r a t u r e s   o f   i 1 3 0 ~ 1 7 6   a n d   o x y g e n   p a r t i a l   p r e s s u r e s   a r o u n d   3 . 9   a n d   2 0   T o r r .   T h e   r a t e   o f   o x i d a t i o n   w a s   m o n i t o r e d   w i t h   a   t h e r m a l   c o n d u c t i v i t y   c e l l .   I n d e p e n d e n t   m e a s u r e m e n t s   w e r e   m a d e   o f   n e t   w e i g h t   g a i n   a n d   q u a n t i t i e s   o f   C O ( g )   a n d   C O s ( g )   e v o l v e d .   O x i   d a t i o n   w a s   s h o w n   t o   b e   n o n p r e f e r e n t i a l ,   i . e . ,   z i r c o n i u m   w a s   o x i d i z e d   a t   t h e   s a m e   r a t e   a s   c a r b o n .   G a s   p h a s e   d i f f u s i o n   c o n t r o l   i m p o s e d   b y   t h e   e x p e r i m e n t a l   s y s t e m   w a s   f r e q u e n t l y   e n c o u n t e r e d .   W h e r e   i t   w a s   p o s s i b l e   t o   o b s e r v e   a   t r u e   c h e m i c a l l y   c o n t r o l l e d   r e a c t i o n   r a t e ,   t h e   k i n e t i c s   a p p e a r e d   t o   b e   l i n e a r .   M i   c r o s c o p i c   e x a m i n a t i o n   o f   t h e   o x i d i z e d   s p e c i m e n s   r e v e a l e d   p r e f e r e n t i a l   o x i d a   t i o n   a l o n g   g r a i n   b o u n d a r i e s .   B e t w e e n   l l 3 0   ~   a n d   1 5 6 0 ~   t h i s   p r e f e r e n t i a l   o x i   d a t i o n   r e s u l t e d   i n   i n t e r c r y s t a l l i n e   f r a c t u r e .   A t   h i g h e r   t e m p e r a t u r e s   s t r e s s e s   w e r e   a p p a r e n t l y   s u f f i c i e n t l y   r e l i e v e d   s o   t h a t   t h e   s a m p l e s   r e m a i n e d   i n t a c t .   T h e   o x i d a t i o n   o f   H f C   b e t w e e n   1 7 9 0   ~   a n d   2 0 0 0 ~   a t   o x y g e n   p r e s s u r e s   n e a r   i 0   T o r r ,   w a s   a l s o   f o u n d   t o   b e   l i n e a r   a n d   p r e f e r e n t i a l   a l o n g   g r a i n   b o u n d a r i e s .   T h e   o x i d a t i o n   o f   z i r c o n i u m   c a r b i d e   w a s   s t u d i e d   b y   M a r g r a v e   a n d   K u r i a k o s e   ( 1 )   a t   t e m p e r a t u r e s   o f   5 5 0   ~   6 5 0 ~   i n   o x y g e n   a t   1   a t m .   T h e   o x i d a t i o n   w a s   f o u n d   t o   b e   l i n e a r ,   w i t h   a n   a c t i v a t i o n   e n e r g y   o f   1 6 . 7   _   1 . 7   k c a l / m o l e .   B a r t l e t t ,   W a d s w o r t h ,   a n d   C u t l e r   ( 2 )   s t u d i e d   t h e   w e i g h t   g a i n   o f   s i z e d   p o w d e r s   o f   z i r c o n i u m   c a r b i d e   i n   a i r ,   o x y g e n ,   a n d   o x y g e n h e l i u m   m i x t u r e s   a t   t e m p e r a   t u r e s   o f   4 5 0 ~ 1 7 6   a n d   o x y g e n   p r e s s u r e s   o f   6 . 5   x   1 0 ~   t o   1   a t m .   S t o i c h i o m e t r i c   o x i d a t i o n   o f   Z r C   t o   Z r O 2   a n d   g a s e o u s   o x i d e s   o f   c a r b o n   w a s   a s s u m e d .   T h e   d a t a   w e r e   i n t e r p r e t e d   o n   t h e   b a s i s   o f   t w o   p a r a l l e l   i n d e   p e n d e n t   p r o c e s s e s :   a   p a r a b o l i c   d i f f u s i o n   r e a c t i o n   i n   v o l v i n g   t h e   p a r t i a l   r e p l a c e m e n t   o f   i n t e r s t i t i a l   c a r b o n   i n   t h e   Z r C   l a t t i c e   w i t h   o x y g e n ,   a n d   a   l i n e a r   s u r f a c e   r e a c t i o n   o c c u r r i n g   a t   t h e   Z r C Z r O 2   p h a s e   b o u n d a r y .   B o t h   p r o c e s s e s   o c c u r   s i m u l t a n e o u s l y ,   w i t h   t h e   d i f f u s i o n   r e a c t i o n   p r e d o m i n a t i n g   a t   s h o r t   t i m e s   a n d   t h e   s u r f a c e   r e a c t i o n   b e c o m i n g   r a t e   c o n t r o l l i n g   a s   o x i d a t i o n   p r o   c e e d s .   T h e   a c t i v a t i o n   e n e r g i e s   w e r e   c a l c u l a t e d   a s   5 3   k c a l / m o l e   f o r   t h e   d i f f u s i o n   p r o c e s s   a n d   4 6   k c a l / m o l e   f o r   t h e   s u r f a c e   r e a c t i o n .   W a t e r   v a p o r ,   i n   t h e   p r e s e n c e   o f   o x y g e n ,   w a s   f o u n d   t o   a c c e l e r a t e   t h e   r a t e   o f   t h e   s u r   f a c e   r e a c t i o n   w h i l e   l e a v i n g   i t s   a c t i v a t i o n   e n e r g y   u n   c h a n g e d .   T h e   s o l i d   o x i d a t i o n   p r o d u c t   w a s   f o u n d   t o   b e   c u b i c   Z r O 2 ,   a   p h a s e   n o r m a l l y   t h o u g h t   t o   b e   u n s t a b l e ,   b u t   w h i c h   m i g h t   b e   s t a b i l i z e d   b y   s m a l l   a m o u n t s   o f   c a r b o n .   W a t t ,   C o c k e t t ,   a n d   H a l l   ( 3 )   m a d e   a   s i n g l e   w e i g h t   c h a n g e   m e a s u r e m e n t   o f   4 9 . 8   m g / c m ~   o n   a   s o l i d   s a m p l e   o f   Z r C   o f   d e n s i t y   6 . 2 0   g / c c   a n d   4 . 8 %   p o r o s i t y ,   e x   p o s e d   t o   a   s t r e a m   o f   d r y   a i r   f l o w i n g   a t   5 . 3   c m / s e c ,   f o r   3 0   m i n   a t   8 0 0 ~   T h e   p r e s e n t   s t u d y   w a s   u n d e r t a k e n   t o   i n v e s t i g a t e   t h e   o x i d a t i o n   o f   Z r C   a n d   t h e   c h e m i c a l l y   r e l a t e d   H f C   a t   t e m p e r a t u r e s   a b o v e   9 0 0 ~   E x p e r i m e n t a l   M e t h o d   C y l i n d r i c a l   p e l l e t s   o f   Z r C   w e r e   c u t   f r o m   z o n e   r e   f i n e d   b a r s   p r e p a r e d   a s   d e s c r i b e d   b y   W e s t r u m   a n d   F e i c k   ( 4 ) .   T h e   f a b r i c a t e d   b a r s   c o n t a i n e d   1 1 . 2   w / o   c a r b o n .   H a f n i u m   c a r b i d e   p o w d e r   w a s   p r e p a r e d   b y   t h e   C a r b o r u n d u m   C o m p a n y   f r o m   h i g h p u r i t y   H f O ~   s u p   p l i e d   b y   W a h   C h a n g   C o r p o r a t i o n .   T h e   H f C   p o w d e r   w a s   s i n t e r e d   i n t o   b a r s   a n d   a r c m e l t e d   o n   a   w a t e r   c o o l e d   c o p p e r   h e a r t h   u s i n g   a   w a t e r c o o l e d   t u n g s t e n   e l e c t r o d e .   I n   o r d e r   t o   m i n i m i z e   l o s s   o f   c a r b o n   d u r i n g   m e l t i n g ,   t h e   o p e r a t i o n   w a s   c o n d u c t e d   i n   a n   a t m o s   p h e r e   o f   a r g o n   c o n t a i n i n g   3 . 1 4 %   o f   e t h y l e n e   a n d   1 1 . 4 %   o f   h y d r o g e n .   T h e   r e s u l t i n g   m a t e r i a l   w a s   c a r b o n   d e   f i c i e n t ,   c o r r e s p o n d i n g   t o   a   c o m p o s i t i o n   H f C 0 . 9 5 2   ( 5 ) .   T h e   a n t i c i p a t e d   p r o d u c t s   o f   t h e   o x i d a t i o n   o f   Z r C   a n d   H f C   w e r e   t h e   p e r m a n e n t   g a s e s ,   C O ( g )   a n d   C O 2 ( g ) ,   i n   a d d i t i o n   t o   t h e   r e f r a c t o r y   m e t a l   o x i d e s .   D u e   t o   t h e   e v o l u t i o n   o f   p e r m a n e n t   g a s e s ,   t h e   t h e r m a l   c o n d u c   t i v i t y   m e t h o d   d e s c r i b e d   i n   p r e v i o u s   p u b l i c a t i o n s   ( 6 8 )   h a d   t o   b e   m o d i f i e d   t o   s t u d y   t h e   o x i d a t i o n .   A   k n o w n   m i x t u r e   o f   h e l i u m   a n d   o x y g e n   w a s   p a s s e d   t h r o u g h   t h e   r e f e r e n c e   s i d e   o f   a   t h e r m a l   c o n d u c t i v i t y   c e l l   ( 6 )   a n d   o v e r   a n   i n d u c t i v e l y   h e a t e d   c a r b i d e   p e l l e t   s u p p o r t e d   o n   T h O 2   f i n g e r s   b y   t h r e e   p o i n t   c o n t a c t .   A   p o r t i o n   o f   t h e   o x y g e n   i n   t h e   g a s   s t r e a m   r e a c t e d   w i t h   t h e   p e l l e t   t o   p r o d u c e   o x i d e s   o f   t h e   m e t a l   a n d   c a r b o n .   T h e   e f f l u e n t   g a s   w a s   t h e r e f o r e   d e p l e t e d   i n   o x y g e n ,   b u t   e n r i c h e d   i n   C O ( g )   a n d   C O 2 ( g ) .   T h e   l a t t e r   w a s   r e m o v e d   b y   p a s s   a g e   t h r o u g h   a   w e i g h e d   A s c a r i t e   b u l b ,   a n d   t h e   r e m a i n   i n g   m i x t u r e   o f   C O ( g ) ,   O 2 ( g ) ,   a n d   H e   e n t e r e d   t h e   s a m p l i n g   s i d e   o f   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   ( 6 ) .   F i n a l l y ,   t h e   C O ( g )   w a s   o x i d i z e d   t o   C O 2 ( g )   o v e r   c o p   p e r   o x i d e   p o w d e r   a t   5 0 0 ~   a n d   t h e   C O s ( g )   p r o d u c e d   w a s   c o l l e c t e d   i n   a   s e c o n d   w e i g h e d   A s c a r i t e   b u l b .   T h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   i n   t h i s   c a s e   w a s   t h e r e f o r e   n o t   d i r e c t l y   r e l a t e d   t o   t h e   r a t e   o f   o x y   g e n   c o n s u m p t i o n   a s   i n   p r e v i o u s l y   s t u d i e d   s y s t e m s   ( 6 8 ) ,   b u t   i n s t e a d   r e f l e c t e d   t h e   d i f f e r e n c e   b e t w e e n   t h e   r a t e   o f   t o t a l   o x y g e n   c o n s u m p t i o n   a n d   t h e   r a t e   o f   f o r m a t i o n   o f   C O ( g ) .   T h e   c a l i b r a t i o n   c o n s t a n t ,   r e   l a t i n g   t h e   e l e c t r i c a l   s i g n a l   t o   t h e   d i f f e r e n c e   i n   g a s   p r e s s u r e   o n   t h e   t w o   s i d e s   o f   t h e   c e l l   i s   t h e   s a m e   f o r   b o t h   C O   a n d   O ~   i n   H e .   H e n c e ,   e v o l u t i o n   o f   C O   d e   p r e s s e s   t h e   s i g n a l ,   a s   c o m p a r e d   t o   s i m p l e   r e m o v a l   o f   o x y g e n .   R e s u l t s   Z i r c o n i u m   c a r b i d e . O x i d a t ~ e n   k i n e t i c s . T h e   e x   p e r i m e n t a l   d a t a   a r e   s u m m a r i z e d   i n   T a b l e   I .   T h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   a p p a r a t u s   w a s   c o n s t a n t   w i t h   t i m e   i n   e v e r y   e x p e r i m e n t .   H o w e v e r ,   o n l y   i n   t h e   c a s e   o f   e x p e r i m e n t   X I I 1   d i d   t h i s   r e f l e c t   t r u e   c h e m i   c a l l y   c o n t r o l l e d   l i n e a r   o x i d a t i o n   k i n e t i c s .   I n   t h e   o t h e r   e x p e r i m e n t s   m o r e   t h a n   9 0 %   o f   t h e   o x y g e n   p a s s e d   o v e r   t h e   c a r b i d e   p e l l e t   r e a c t e d   w i t h   i t ,   a n d   t h e   c o n t r o l l i n g   p r o c e s s   w a s   t h e r e f o r e   p r o b a b l y   t h e   r a t e   o f   a r r i v a l   o f   o x y g e n   g a s   a t   t h e   s a m p l e   s u r f a c e .   A t   h i g h e r   p r e s s u r e s   a n d / o r   h i g h e r   g a s   f l o w   r a t e s ,   a   g r e a t e r   p r o p o r t i o n   o f   t h e   c a r b i d e   w o u l d   h a v e   b e e n   c o n v e r t e d   t o   o x i d e s .   I n   T a b l e   I ,   t h e   \" i n i t i a l \"   w e i g h t s   w e r e   t a k e n   a f t e r   d e g a s s i n g   a t   2 2 0 0 ~   i n   p u r e   h e l i u m   u n t i l   t h e   s i g n a l   f r o m   t h e   t h e r m a l   c o n d u c t i v i t y   c e l l   i n d i c a t e d   t h a t   n o   p e r m a n e n t   g a s e s   w e r e   b e i n g   e v o l v e d .   T h e   s u r f a c e   a r e a s   w e r e   c a l c u l a t e d   f r o m   m i c r o m e t e r   m e a s u r e m e n t s   o f   t h e   h e i g h t   a n d   d i a m e t e r   o f   t h e   c y l i n d r i c a l   p e l l e t s .   T e m p e r a t u r e s   w e r e   m e a s u r e d   o p t i c a l l y   a n d   c o r r e c t e d   f o r   a n   e m i s s i v i t y   o f   0 . 7 ,   d e t e r m i n e d   b y   c o m p a r i n g   1 0 3 0   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   V o l .   1 1 4 ,   N o .   1 0   H I G H T E M P E R A T U R E   O X I D A T I O N   T a b l e   I .   S u m m a r y   o f   e x p e r i m e n t a l   d a t a   o n   Z r C   1 0 3 1   G e o m e t r i c   I n i t i a l   s u r f .   a r e a ,   E x p t .   w e i g h t ,   g   c m   2   T e m p ,   * I 4 [   C O z   O x y g e n   E x p o s u r e   N e t   w e i g h t   C O   f o r m e d   C   Z r   p r e s s u r e ,   t i m e ,   c h a n g e   f o r m e d   W c o   ( g )   c o n s u m e d   c o n s u m e d   T o r r   r a i n   W o ,   ( g )   W o o   ( g )   W e   ( g )   W z r   ( g )   Z r / C   X I I 8   0 . 4 4 6   1 . 1 0 3   1 1 3 0   X I I 5   0 . 5 6 4 5   1 . 3 9 4   1 2 6 0   X I I 3   0 . 6 9 6 3   1 . 4 9 3   1 5 6 0   X 3 1   0 . 5 8 6 4   1 . 3 1 0   1 8 6 0   X 2 9   0 . 7 0 7 3   1 . 4 5 0   1 9 4 0   X I I I   0 . 7 0 5 1   1 . 5 0 0   1 9 7 0   X 2 7   0 . 6 3 6 1   1 . 2 5 2   2 0 7 0   X 1 6   0 . 6 8 4 0   1 . 3 8 6   2 0 7 0   V I I 8   0 . 6 6 4 3   1 . 5 2 8   2 1 0 0   X 2 5   0 . 6 8 1 2   1 . 3 7 1   2 1 6 5   2 2 . 9   5 1     0 . 0 0 0 1   0 . 0 6 1 1         2 0 . 4   1 2 8       0 . 0 7 0 9     ~     2 1 . 2   6 2     0 . 0 2 8 9   0 . 0 3 9 9       9 . 1   1 2 9   0 . 0 4 0 2   0 . 0 4 3 8   0 . 0 2 2 0   0 . 0 2 4 8   0 . 1 8 5 2   8 . 1   1 2 0   0 . 0 3 6 6   0 . 0 4 0 5   0 . 0 1 4 5   0 . 0 2 1 4   0 . 1 6 5 2   7 . 7 2   2 5 . 9   1 1 2   0 . 0 8 5 1   0 . 0 9 6 8   0 . 0 3 5 6   0 . 0 5 1 0   0 . 3 8 8   7 . 6 0   8 . 5   1 2 4   0 . 0 3 8 9   0 . 0 4 6 4   0 . 0 1 7 5   0 . 0 2 3 6   0 . 1 7 8 0   7 . 5 5   6 . 5   1 1 9   0 . 0 4 0 2   0 . 0 5 4 0   0 . 0 0 7 3   0 . 0 2 5 2   0 . 1 8 6 2   7 . 3 9   3 . 0   1 8 0   0 . 0 2 0 7   0 . 0 2 4 2   0 . 0 0 7 9   0 . 0 1 2 6   0 . 0 9 4 9   7 . 5 4   8 . 9   1 2 0   0 . 0 3 5 6   0 . 0 5 2 3   0 . 0 0 7 2   0 . 0 2 4 4   0 . 1 7 1 0   7 . 0 1   s u r f a c e   t e m p e r a t u r e s   w i t h   u l t r a s o n i c a l l y   d r i l l e d   b l a c k   b o d y   c a v i t y   t e m p e r a t u r e s   u n d e r   o x i d i z i n g   c o n d i t i o n s .   O x y g e n   p a r t i a l   p r e s s u r e s   i n   h e l i u m   a r e   g i v e n ;   t h e   t o t a l   p r e s s u r e   w a s   c l o s e   t o   1   a t m   i n   e v e r y   e x p e r i m e n t .   T h e   c a r r i e r   g a s   f l o w   r a t e   w a s   5 8 . 6   c c / m i n   i n   e v e r Y   e x   p e r i m e n t ,   c o r r e s p o n d i n g   t o   a   l i n e a r   f l o w   v e l o c i t y   i n   t h e   n e i g h b o r h o o d   o f   t h e   s a m p l e   o f   1 . 9   c m / s e c ,   e x c e p t   V I I 8   w h e r e   a   f l o w   o f   5 1 . 5   c c / m i n   w a s   u s e d .   F r o m   t h e   m e a s u r e d   n e t   w e i g h t   c h a n g e   o f   t h e   c a r   b i d e   o n   o x i d a t i o n ,   a n d   t h e   o b s e r v e d   w e i g h t   c h a n g e s   W c o 2   a n d   W c o   i n   t h e   A s c a r i t e   b u l b s ,   t h e   t o t a l   a m o u n t s   o f   c a r b o n   a n d   z i r c o n i u m   c o n s u m e d   w e r e   r e a d i l y   c a l   c u l a t e d ,   o n   t h e   a s s u m p t i o n   t h a t   t h e   o n l y   o x i d a t i o n   p r o d u c t s   w e r e   C O 2 ( g ) ,   C O ( g ) ,   a n d   Z r O ~ ( s ) .   T h e   f o r m a t i o n   o f   Z r O C   ( s )   c a n n o t   b e   p r e c l u d e d ,   b u t   s i n c e   i t   i s   i s o s t r u c t u r a l   w i t h   Z r C ,   n o   p o s i t i v e   e v i d e n c e   w a s   o b   t a i n e d   f o r   i t s   p r e s e n c e .   T h e   t o t a l   w e i g h t   o f   c a r b o n   c o n s u m e d ,   W e ,   i s   g i v e n   b y   [ c ]   [ c ]   W e   =     (cid:12) 9   W c o 2   +   ~   \" W c o   [ 1 ]   [ C O ~ ]   [ C O ]   w h e r e   t h e   s y m b o l s   i n   b r a c k e t s   r e p r e s e n t   m o l e c u l a r   w e i g h t s .   T h e   t o t a l   w e i g h t   o f   z i r c o n i u m ,   W z r ,   t h a t   h a s   b e e n   c o n v e r t e d   t o   o x i d e   i s   c a l c u l a t e d   f r o m   t h e   m e a   s u r e d   w e i g h t   c h a n g e ,   W o ,   a n d   t h e   d e r i v e d   c a r b o n   c o n   s u m p t i o n   [ Z r ]   W z r   =   ~   ( W o   +   W c )   [ 2 ]   2 [ 0 ]   T h e   r a t i o   o f   t h e   n u m b e r   o f   g r a m s   o f   z i r c o n i u m   c o n   s u m e d   t o   t h e   n u m b e r   o f   g r a m s   o f   c a r b o n   c o n s u m e d   d u r i n g   o x i d a t i o n   i s   s h o w n   i n   T a b l e   I   t o   h a v e   a n   a p   p r o x i m a t e l y   c o n s t a n t   v a l u e   o f   7 . 5   (cid:127)   0 . 2 .   S i n c e   t h e   c o r r e s p o n d i n g   r a t i o   i n   t h e   Z r C   s t a r t i n g   m a t e r i a l   i s   7 . 6 ,   i t   w o u l d   a p p e a r   t h a t   t h e   o x i d a t i o n   o f   Z r C   i s   s t o i   c h i o m e t r i c .   T h a t   i s ,   f o r   e a c h   z i r c o n i u m   a t o m   c o n v e r t e d   t o   o x i d e ,   a   s i n g l e   c a r b o n   a t o m   i s   a l s o   c o n v e r t e d   t o   o x i d e .   S t r u c t u r a l   c h a n g e s   d u r i n g   o x i d a t i o n . T h e   r e a s o n   t h a t   w e i g h t   c h a n g e   d a t a   w e r e   n o t   g i v e n   f o r   p e l l e t s   X I I 8 ,   X I I 5 ,   a n d   X I I 3   i s   t h a t   a t   t h e s e   r e l a t i v e l y   l o w   t e m p e r a t u r e s   t h e   p e l l e t s   w e r e   b r o k e n   a p a r t   b y   t h e   o x i d a t i o n   p r o c e s s .   A t   t h e   e n d   o f   e a c h   e x p e r i m e n t ,   t h e   g r a i n   b o u n d a r i e s   o f   t h e   Z r C   w e r e   s e e n   t o   b e   o u t l i n e d   b y   a   w h i t e   m a t e r i a l ,   i d e n t i f i e d   b y   r o o m   t e m p e r a t u r e   x r a y   a s   m o n o c l i n i c   Z r O 2 .   T h e   g r o w t h   o f   t h e   o x i d e   i n   p r e e x i s t i n g   c r a c k s   a n d   g r a i n   b o u n d a r i e s   o f   Z r C   u n   d o u b t e d l y   c r e a t e s   e n o u g h   s t r e s s   t o   f r a c t u r e   t h e   c a r   b i d e .   A t   h i g h e r   t e m p e r a t u r e s ,   a s   d i s c u s s e d   b e l o w ,   s i g   n i f i c a n t   g r a i n   b o u n d a r y   o x i d a t i o n   w a s   o b s e r v e d ,   b u t   s t r e s s e s   a r e   a p p a r e n t l y   s u f f i c i e n t l y   r e l i e v e d   s o   t h a t   f r a c t u r e   d o e s   n o t   o c c u r .   P e l l e t   X I I 1   f o r   w h i c h   t r u e   k i n e t i c   d a t a   h a d   b e e n   o b t a i n e d   w a s   m o u n t e d   a n d   p o l i s h e d   f o r   m e t a l l o g r a p h i c   e x a m i n a t i o n .   F i g u r e   l a   s h o w s   t h e   s p e c i m e n   a t   a   m a g n i f i c a t i o n   o f   c a .   4 X .   T h e   g r a y   o u t e r   r i m   i s   t h e   o x i d e ,   a n d   t h e   i n n e r   w h i t e   c i r c u l a r   a r e a   i s   t h e   s u r f a c e   o f   c a r b i d e .   T o   t h e   n a k e d   e y e ,   t h e   o u t e r   o x i d e   c o a t i n g   l o o k e d   w h i t e   a n d   c h a l k y ,   a n d   t h e   i n n e r   s u r f a c e   f r o m   w h i c h   t h e   o x i d e   c o a t i n g   w a s   p o l i s h e d   o f f ,   l o o k e d   b r i g h t   a n d   m e t a l l i c .   A   m o t t l e d   r i m   i s   c l e a r l y   v i s i b l e   a l o n g   t h e   o x i d e a l l o y   i n t e r f a c e   i n   t h e   f i g u r e   a n d   t h e   r e m a i n i n g   p h o t o m i c r o   g r a p h s   f o c u s   o n   p o r t i o n s   o f   t h i s   i n t e r f a c e .   I n   F i g .   l b ,   a t   a   m a g n i f i c a t i o n   o f   c a .   3 0 X ,   t h e   o x i d e   f i l l s   t h e   e n t i r e   u p p e r   h a l f   o f   t h e   p h o t o g r a p h ,   a n d   t h e   a l l o y   t h e   l o w e r .   T h e   o x i d e   i s   o b v i o u s l y   g r o w i n g   p r e f e r e n t i a l l y   a l o n g   g r a i n   b o u n d a r i e s   i n   t h e   c a r b i d e ,   a n d   e n v e l o p i n g   i n   d i v i d u a l   a l l o y   c r y s t a l l i t e s .   T h e   s t r u c t u r e   o f   t h e   b u l k   o x i d e   i s   v e r y   d i f f e r e n t   f r o m   t h a t   o n   Z r B 2   ( 8 ) ,   a l t h o u g h   a f t e r   c o o l i n g   b o t h   s h o w e d   o n l y   t h e   x r a y   l i n e s   f o r   m o n o c l i n i c   Z r O 2 .   O n   Z r B 2   t h e   o x i d e   w a s   s e e n   t o   g r o w   i n   a   c o l u m n a r   s t r u c t u r e ;   o n   Z r C ,   t h e   Z r O 2   a s s u m e d   a n   e q u i a x e d   g r a i n   s t r u c t u r e   v e r y   s i m i l a r   t o   t h a t   o f   t h e   o r i g i n a l   a l l o y .   F i g u r e s   l c   a n d   l d   s h o w   t h e   o x i d e   a l l o y   i n t e r f a c e   a t   a   s t i l l   h i g h e r   m a g n i f i c a t i o n ,   c a .   1 9 5 X ,   a n d   o n e   s e e s   e v e n   m o r e   c l e a r l y   t h e   p r e f e r e n t i a l   o x i d a t i o n   o f   g r a i n   b o u n d a r i e s ,   a n d   t h e   l a t e r a l   f i n g e r l i k e   g r o w t h   o f   o x i d e   f r o m   t h e   b o u n d a r i e s   i n t o   t h e   c r y s t a l l i t e   b u l k .   T h e   r e a c t i o n   z o n e   o f   i n t e r g r a n u l a r   a t t a c k   w a s   a p   p r o x i m a t e l y   0 . 0 1 4   + _   0 . 0 0 2   c m   i n   t h i c k n e s s .   T h e   o u t e r   o x i d e   w a s   a b o u t   t e n   t i m e s   t h i c k e r .   T h e   m e c h a n i s m   o f   o x i d e   g r o w t h   o n   Z r C   a t   h i g h   t e m p e r a t u r e s   a p p e a r s   t o   b e   r a p i d   a t t a c k   a t   g r a i n   b o u n d a r i e s ,   a n d   s l o w   o x i d a   t i o n   o f   t h e   a l l o y   f r o m   t h e   g r a i n   b o u n d a r y   s u r f a c e   i n   w a r d .   F i g u r e   l e ,   a t   a   m a g n i f i c a t i o n   o f   c a .   3 4 5 X ,   s h o w s   p o r t i o n s   o f   t h e   a l l o y   c o m p l e t e l y   e n v e l o p e d   b y   o x i d e .   F i n a l l y ,   F i g .   I f   i s   a   v i e w   o f   t h e   a l l o y   s u r f a c e   a t   a   m a g   n i f i c a t i o n   o f   c a .   3 4 5 X .   T h i s   s u r f a c e   h a d   b e e n   c o v e r e d   w i t h   a   d e n s e   o x i d e   p r i o r   t o   p o l i s h i n g ,   a n d   o n e   s e e s   h e r e   t h e   p e n e t r a t i o n   o f   t h e   o x i d e   i n t o   g r a i n   b o u n d a r i e s   o f   t h e   a l l o y .   B e t w e e n   1 1 2 6   ~   a n d   1 5 5 9 ~   t h e   g r a i n   b o u n d a r y   a t t a c k   r e s u l t s   i n   i n t e r g r a n u l a r   f r a c t u r e   o f   t h e   a l l o y .   A b o v e   1 5 8 0 ~   t h e r e   i s   a p p a r e n t l y   e n o u g h   p l a s t i c i t y   i n   e i t h e r   a l l o y ,   o x i d e ,   o r   b o t h ,   s o   t h a t   t h e   s a m p l e   r e m a i n s   i n t a c t   d u r i n g   o x i d a t i o n .   H a f n i u m   c a r b i d e . O x i d a t i o ~   k i n e t i c s . T h e   e x p e r i   m e n t a l   d a t a   f o r   H f C   a r e   s u m m a r i z e d   i n   T a b l e   I I .   S i n c e   t h e   a r c m e l t e d   s a m p l e s   w e r e   h i g h l y   i r r e g u l a r   i n   s h a p e ,   F i g .   1   a d .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   Z r C   ( p e l l e t   X I I 1 ) ;   T   1 9 6 ~   P o 2   2 5 . 9   T a r r .   M a g n i f i c a t i o n :   a   ( t a p   l e f t )   c a .   4 X ;   b   ( t a p   r i g h t )   c a .   3 0 X ;   c   ( b o t t o m   l e f t )   c a .   1 9 5 X ;   d   ( b o t t o m   r i g h t )   c a .   1 9 5 X .   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   1 0 3 2   J .   E l e c t r o c h e m .   S o c . :   E L E C T R O C H E M I C A L   S C I E N C E   O c t o b e r   1 9 6 7   F i g .   1   e f .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   Z r C   ( p e l l e t   X I I 1 ) ;   T   1 9 6 ~   P o 2   2 5 . 9   T o r r .   M a g n i f i c a t i o n :   e   ( l e f t )   a n d   f   ( r i g h t )   c a .   3 4 5 X .   F i g .   2 .   P h o t o m i c r o g r a p h s   o f   o x i d i z e d   h a f n i u m   c a r b i d e   p e l l e t   X V I 1 9 ;   t e m p e r a t u r e   1 6 0 0 ~ 1 7 6   P o 2   1 1 . 5   T o r r .   M a g n i f i c a t i o n   c a .   1 3 0 X .   t h e   g e o m e t r i c a l l y   c a l c u l a t e d   s u r f a c e   a r e a s   a r e   o n l y   a p p r o x i m a t e .   T h e   s u r f a c e   o x i d e   t e n d e d   t o   f l a k e   a n d   s p a l l   o n   r e m o v a l   o f   t h e   s a m p l e   f r o m   t h e   a p p a r a t u s ;   a n d ,   h e n c e ,   t h e   m e a s u r e d   w e i g h t   g a i n s   a r e   m i n i m u m   v a l u e s .   F r o m   t h e   m e a s u r e d   m i n i m u m   n e t   w e i g h t   g a i n s ,   a n d   t h e   w e i g h t s   o f   C O ( g )   a n d   C O s ( g )   i n   t h e   p r o d u c t   g a s   s t r e a m ,   t h e   m i n i m u m   r a t i o   o f   h a f n i u m   t o   c a r b o n   o x i d i z e d   c a n   b e   c a l c u l a t e d ,   a s   d e s c r i b e d   a b o v e   f o r   t h e   o x i d a t i o n   o f   Z r C .   T h e   e x p e r i m e n t a l   v a l u e s   o f   t h e   r a t i o   a r e   g i v e n   i n   T a b l e   I I   a s   1 4 . 1 ,   1 0 . 1 ,   a n d   1 5 . 9 ,   t o   b e   c o m   p a r e d   t o   a n   H f : C   r a t i o   i n   t h e   o r i g i n a l   a l l o y   o f   1 5 . 6 .   L i n e a r   o x i d a t i o n   w a s   o b s e r v e d   i n   e v e r y   c a s e .   F l o w   r a t e s   o f   1 1 9   c c / m i n ,   a p p r o x i m a t e l y   d o u b l e   t h o s e   u s e d   f o r   t h e   Z r C   e x p e r i m e n t s ,   w e r e   e m p l o y e d   i n   o r d e r   t o   m a i n t a i n   a n   a d e q u a t e   s u p p l y   o f   o x y g e n .   A t   1 2 8 0 ~   i n   p u r e   o x y g e n   a t   1   a t m ,   a   h a f n i u m   c a r   b i d e   s p e c i m e n   d i s i n t e g r a t e d   i n t o   s e v e r a l   p i e c e s   w i t h i n   3   m i n   i n   a   m a n n e r   v e r y   s i m i l a r   t o   t h a t   d e s c r i b e d   a b o v e   f o r   z i r c o n i u m   c a r b i d e .   M e t a l l o g r a p h i c   e x a m i n a t i o n   o f   o x i d e   f i l m s . ~ P e l l e t   X V I 1 9   w a s   i m b e d d e d   i n   p l a s t i c   a n d   p o l i s h e d   f o r   m i   c r o s c o p i c   o b s e r v a t i o n .   I n   t h e   p h o t o m i c r o g r a p h s   s h o w n   i n   F i g .   2 ,   t h e   o x i d e   i s   s e e n   t o   c o n t a i n   c o n s i d e r a b l e   p o   r o s i t y .   A s   i n   t h e   c a s e   o f   z i r c o n i u m   c a r b i d e ,   o x i d a t i o n   a p p e a r s   t o   b e   p r e f e r e n t i a l   a l o n g   g r a i n   b o u n d a r i e s .   T h e   p h o t o m i c r o g r a p h   i n   F i g .   3   i s   f o r   a   h a f n i u m   c a r   b i d e   p e l l e t   t h a t   h a d   b e e n   e x p o s e d   t o   a   N 2 H e   m i x t u r e   f o r   a n   h o u r   a n d   a   q u a r t e r   a t   a   t e m p e r a t u r e   o f   I 9 6 0 ~   T h e   N u   a p p a r e n t l y   c o n t a i n e d   a   s m a l l   q u a n t i t y   o f   o x y   g e n ,   a n d   w h i t e   H f O 2 ( c )   w a s   i d e n t i f i e d   o n   t h e   s u r f a c e   o f   t h e   s a m p l e   b y   x r a y   d i f f r a c t i o n .   T h e   n e t   w e i g h t   c h a n g e   o f   t h e   s a m p l e   w a s   0 . 0 0 2 4   g / c m   2 ,   a n   o r d e r   o f   m a g n i t u d e   l e s s   t h a n   t h a t   o b s e r v e d   a t   a n   o x y g e n   p a r   t i a l   p r e s s u r e   o f   1 1 . 5   T o r r   a t   a b o u t   t h e   s a m e   t e m p e r a   t u r e .   T h e   o x i d a t i o n   r a t e   w a s   s e e n   t o   d e c r e a s e   s l i g h t l y   w i t h   t i m e ,   a n d ,   a l t h o u g h   g r a i n   b o u n d a r y   o x i d a t i o n   i s   T a b l e   I I .   S u m m a r y   o f   F i g .   3 .   P h o t o m i c r o g r a p h   a f   o x i d i z e d   h a f n i u m   c a r b i d e   p e l l e t   X V I 2 2 ;   t e m p e r a t u r e   1 9 6 0 ~   l o w   o x y g e n   p r e s s u r e .   M a g n i f i c a t i o n   c a .   2 3 0 X .   a p p a r e n t   i n   F i g .   3 ,   t h e   o x i d e   s e e m s   l e s s   p o r o u s   t h a n   t h a t   s h o w n   i n   F i g .   2 ,   f o r   o x i d a t i o n   a t   a   h i g h e r   p r e s   s u r e .   D i s c u s s i o n   A   r e c e n t   r e p o r t   ( 9 )   s u g g e s t s   t h a t   g r a i n   b o u n d a r y   o x i d a t i o n   i s   c h a r a c t e r i s t i c   O f   i r o n   c o n t a m i n a t e d   Z I C .   F o r   m a t e r i a l   w i t h   a n   i m p u r i t y   l e v e l   a b o u t   1   v / o   F e s C   o r   a b o u t   1 . 1   w / o ,   t h e   F e 3 C   i s   f o u n d   t o   b e   s e g r e g a t e d   a t   t h e   g r a i n   b o u n d a r i e s ,   a n d   t o   b e   o x i d i z e d   a t   a   m u c h   m o r e   r a p i d   r a t e   t h a n   t h e   Z r C   m a t r i x .   A   c h e m i c a l   a n a l y s i s   o f   t h e   z o n e   m e l t e d   m a t e r i a l   u s e d   i n   t h e   p r e s e n t   s t u d y   s h o w e d   a n   i r o n   c o n t a m i n a   t i o n   l e v e l   o f   0 . 0 7 %   b y   w e i g h t .   I n   o r d e r   t o   t r y   t o   a s s e s s   t h e   p o s s i b l e   i n f l u e n c e   o f   i r o n   o n   t h e   p r e s e n t   r e s u l t s ,   t h e   o x i d i z e d   a n d   p o l i s h e d   Z r C   s a m p l e   s h o w n   i n   F i g .   1   w a s   e x a m i n e d   f o r   i r o n   w i t h   t h e   e l e c t r o n   p r o b e .   1   A n   F e K ~ I   s c a n n i n g   i m a g e   w a s   t a k e n   o f   t h e   s a m p l e ,   a n d   i r o n   w a s   f o u n d   t o   b e   c o n c e n t r a t e d   i n   t h e   s m a l l   c i r c u l a r   i n c l u s i o n s   v i s i b l e   i n   F i g .   1 ,   a n d   i n h o m o g e n e   o u s l y   i n   t h e   c r a c k s .   P o i n t   c o u n t   a n a l y s e s   t a k e   i n   t h e   i n c l u s i o n s ,   t h e   c r a c k s ,   t h e   g r a i n   b o u n d a r i e s ,   a n d   t h e   m a t r i x   s h o w e d   t h e   p r e s e n c e   o f   i r o n   p a r t i c l e s   o f   l e s s   t h a n   1 ~   i n   s i z e   d i s t r i b u t e d   a t   r a n d o m   i n   t h e   i n c l u s i o n s   a n d   c r a c k s ,   b u t   f a i l e d   t o   r e v e a l   t h e   p r e s e n c e   o f   a n y   i r o n   a t   a l l   i n   e i t h e r   t h e   g r a i n   b o u n d a r i e s   o r   t h e   m a t r i x .   T h e   p r e f e r e n t i a l   g r a i n   b o u n d a r y   o x i d a t i o n   o b s e r v e d   i n   z i r c o n i u m   c a r b i d e ,   t h e r e f o r e ,   s e e m s   t o   b e   c h a r a c t e r   i s t i c   o f   t h e   p u r e   m a t e r i a l .   A c k n o w l e d g m e n t   I t   i s   a   p l e a s u r e   t o   t h a n k   J o h n   E n g e l k e   o f   A r t h u r   D .   L i t t l e ,   I n c . ,   f o r   h i s   a s s i s t a n c e   i n   o b t a i n i n g   a n d   i n t e r   p r e t i n g   t h e   e l e c t r o n   p r o b e   d a t a ,   a n d   t o   a c k n o w l e d g e   t h e   i n v a l u a b l e   a s s i s t a n c e   o f   R i c h a r d   F .   Q u i g l e y   a n d   W a l t e r   C h r i s t e n s e n   w i t h   t h e   e x p e r i m e n t a l   w o r k .   M a n u s c r i p t   r e c e i v e d   J a n .   2 3 ,   1 9 6 7 ;   r e v i s e d   m a n u   s c r i p t   r e c e i v e d   M a y   3 1 ,   1 9 6 7 .   A n y   d i s c u s s i o n   o f   t h i s   p a p e r   w i l l   a p p e a r   i n   a   D i s c u s s i o n   S e c t i o n   t o   b e   p u b l i s h e d   i n   t h e   J u n e   1 9 6 8   J O U R N A L .   1 T h e   e l e c t r o n   p r o b e   a n a l y s i s   w a s   k i n d l y   s u p p l i e d   b y   A c t o n   L a b o r a t o r i e s ,   5 3 1   M a i n   S t r e e t ,   A c t o n ,   M a s s a c h u s e t t s .   e x p e r i m e n t a l   d a t a   o n   H f C   I n i t i a l   P r e s s u r e ,   T i m e ,   E x p t .   w e i g h t ,   g   A r e a ,   c m   ~   T e m p ,   \" K   T o r r   m i n   M i n i m u m   M i n i m u m   C   H f   n e t   w e i g h t   C O   C O s   c o n s u m e d ,   c o n s u m e d ,   H f / C   c h a n g e ,   g   f o r m e d ,   g   f o r m e d ,   g   g   g   r a i n   X V I I 4   0 . 5 3 7 4   1 , 0 7 6   1 7 9 0   1 1 . 5   8 5   X V I 3 8   0 . 6 5 4 4   1 . 2 4 1   1 8 9 0   1 1 . 5   8 9   X V I 3 6   0 . 7 5 1 1   1 . 0 3 2   2 0 0 5   1 1 . 5   4 3   X V I 1 9   0 . 6 0 5 7   1 . 2 5 6   2 0 0 0   7 . 5   1 2 0   1 6 0 0   ( c o n t i n u o u s   d r o p )   0 . 0 2 1   0 . 0 1 6 0   0 . 0 2 5 4   0 . 0 1 3 8   0 . 1 9 5   1 4 . 1   0 . 0 1 6   0 . 0 2 3 2   0 . 0 3 5 5   0 . 0 1 9 6   0 . 1 9 9   1 0 . 1   0 . 0 2 2     0 . 0 1 9 1   0 , 0 3 8   0 . 0 3 4 0   0 . 0 2 2 5   0 . 0 2 0 7   0 , 3 2 8   1 5 . 9   \\x0c', 'Downloaded on 2014-11-01 to IP   155.247.166.234   address. Redistribution subject to ECS terms of use (see   ecsdl.org/site/terms_use  ) unless CC License in place (see abstract).(cid:160)   V o l .   1 1 4 ,   N o .   1 0   H I G H T E M P E R A T U R E   O X I D A T I O N   1 0 3 3   R E F E R E N C E S   1 .   A .   K .   K u r i a k o s e   a n d   J .   L .   M a r g r a v e ,   T h i s   J o u r n a l ,   1 1 1 ,   8 2 7   ( 1 9 6 ~ t ) .   2 .   R .   W .   B a r t l e t t ,   M .   E .   W a d s w o r t h ,   a n d   I .   B .   C u t   l e r ,   T r a n s .   A I M E ,   2 2 7 ,   4 6 7   ( 1 9 6 3 ) .   3 .   W .   W a t t ,   G .   H .   C o c k e t t ,   a n d   A .   R .   H a l l ,   M e t a u x ,   2 8 ,   2 2 2   ( 1 9 6 3 ) .   4 .   E .   F .   W e s t r u m ,   J r . ,   a n d   G .   F e i c k ,   J .   C h e m .   E n g .   D a t a ,   8 ,   1 7 6   ( 1 9 6 3 ) .   5 .   L .   A .   M c C l a i n e ,   T h e r m o d y n a m i c   a n d   K i n e t i c   S t u d i e s   f o r   a   R e f r a c t o r y   M a t e r i a l s   P r o g r a m ,   4 t h   S e m i   a n n u a l   P r o g r e s s   R e p o r t   ( A u g u s t ,   1 9 6 3 ) ,   C o n t r a c t   N o .   A F   3 3   ( 6 1 6 ) 7 4 7 2 .   6 .   J .   B .   B e r k o w i t z N I a t t u c k ,   T h i s   J o u r n a l ,   I l l ,   9 0 8   ( 1 9 6 4 ) .   7 .   J .   B .   B e r k o w i t z M a t t u c k   a n d   R .   R .   D i l s ,   i b i d . ,   1 1 2 ,   5 8 3   ( 1 9 6 5 ) .   8 .   J .   B .   B e r k o w i t z M a t t u c k ,   i b i d . ,   1 1 3 ,   9 0 8   ( 1 9 6 6 ) .   9 .   K .   R .   J a n o w s k i ,   R .   D .   C a r n a h a n ,   a n d   R .   C .   R o s s i ,   \" S t a t i c   a n d   D y n a m i c   O x i d a t i o n   o f   Z r C , \"   T D R   6 6 9   ( 6 2 5 0 1 0 ) 3 ,   A e r o s p a c e   C o . ,   E l   S e g u n d o ,   C a l i f . ,   J a n u a r y ,   1 9 6 6 .   F u n d a m e n t a l   L i m i t a t i o n s   o n   t h e   L o w T e m p e r a t u r e   O p e r a t i o n   o f   E l e c t r o l y t i c   D e v i c e s   C .   A .   A n g e l l   D e p a r t m e n t   o f   C h e m i s t r y ,   P u r ~ u e   U n i v e r s i t y ,   L a f a y e t t e ,   I n d i a n a   A B S T R A C T   A   n e w   a p p r o a c h   t o   t h e   u n d e r s t a n d i n g   o f   t r a n s p o r t   b e h a v i o r   i n   c o n c e n   t r a t e d   e l e c t r o l y t e   s o l u t i o n s   r e q u i r e s   t h e   r e c o g n i t i o n   a t   l o w   t e m p e r a t u r e s   o f   a   l i q u i d   s t a t e   l i m i t i n g   t e m p e r a t u r e   w h i c h   i s   a   t h e r m o d y n a m i c   c o n s t a n t   o f   a n y   s o l v e n t e l e c t r o l y t e   s o l u t i o n .   T h e   v a l u e   o f   t h i s   c o n s t a n t ,   w h i c h   d e t e r m i n e s   s e r v i c e   l i m i t s ,   i s   a   f u n c t i o n   o f   e l e c t r o l y t e   c o n c e n t r a t i o n   e x c e p t   a t   l o w   c o n   c e n t r a t i o n s   w h e r e   s o l v e n t   s t r u c t u r e   m a y   d o m i n a t e .   I n   t h i s   p a p e r   t h e   r e s u l t s   o f   w o r k e r s   i n v e s t i g a t i n g   e l e c t r o l y t e   s o l u t i o n s   f o r   l o w t e m p e r a t u r e   b a t t e r y   a p p l i c a t i o n s   a r e   u s e d   t o   e m p h a s i z e   t h e   u s e f u l n e s s   o f   t h e   c o n c e p t s   u n d e r l y i n g   t h i s   a p p r o a c h   t o   l o w t e m p e r a t u r e   e l e c t r o l y t e   p r o b l e m s .   T h i s   a r t i c l e   c o n s i d e r s   b r i e f l y   s o m e   i m p l i c a t i o n s   o f   r e c e n t   d e v e l o p m e n t s   i n   e l e c t r o l y t e   s o l u t i o n   t r a n s p o r t   t h e o r y   t o   t h e   p r a c t i c a l   p r o b l e m s   e n c o u n t e r e d   i n   t h e   u s e   o f   e l e c t r o l y t e   s o l u t i o n s   a t   l o w   t e m p e r a t u r e s .   T e m p e r a t u r e   a n d   c o n c e n t r a t i o n   d e p e n d e n c e   o f   s o t u   t i o n   t r a n s p o r t   p r o p e r t i e s . I n   r e c e n t   p a p e r s   ( 1 3 )   i t   h a s   b e e n   s h o w n   t h a t   t h e   t e m p e r a t u r e   d e p e n d e n c e   a n d   c o m p o s i t i o n   d e p e n d e n c e   o f   e l e c t r i c a l   ( e q u i v a l e n t )   c o n d u c t a n c e ,   A ,   a n d   v i s c o u s   f l o w ,   ~ ] ,   p r o c e s s e s   i n   v a r i   o u s   c o n c e n t r a t e d   a q u e o u s   e l e c t r o l y t e   s o l u t i o n s   a t   l o w   c o r r e s p o n d i n g   t e m p e r a t u r e s   m a y   b e   d e s c r i b e d   b y   e q u a t i o n s   o f   t h e   f o l l o w i n g   f o r m   c o n s t a n t   c o m p o s i t i o n   k   A ( I ~ )   ,   1 / ~ ] ( N I   =   A   e x p   .   T     T o   c o n s t a n t   t e m p e r a t u r e   k / Q   A ( T )   ,   1 / ~ ] ( T )   . ~   A   e x p   N o     N   [ 1 ]   [ 2 ]   w h e r e   A ,   k ,   a n d   T o   i n   E q .   [ 1 ]   a n d   A ,   k / Q ,   a n d   N o   i n   E q .   [ 2 ]   a r e   c o n s t a n t s ;   T   a n d   N   a r e ,   r e s p e c t i v e l y ,   t h e   a b s o l u t e   t e m p e r a t u r e   a n d   t h e   e q u i v a l e n t   c o n c e n t r a t i o n .   A   p l a u s i b l e   t h e o r e t i c a l   i n t e r p r e t a t i o n   d i s c u s s e d   i n   d e t a i l   e l s e w h e r e   ( 1 ,   4 ,   5 )   m a y   b e   g i v e n   t h e s e   e q u a t i o n s   i n   t e r m s   o f   t h e   c o n t r o l l i n g   i n f l u e n c e   o n   t h e   t r a n s p o r t   p r o c e s s e s ,   o f   t h e   l i q u i d   c o n f i g u r a t i o n a l   e n t r o p y   c o n   t e n t   ( 6 ) .   T h e   c o n s t a n t s   T o   a n d   N o   i n   t h i s   c a s e   r e p r e   s e n t   t h e   t e m p e r a t u r e   a t   f i x e d   c o m p o s i t i o n ,   a n d   t h e   e q u i v a l e n t   c o n c e n t r a t i o n   a t   c o n s t a n t   t e m p e r a t u r e ,   r e   s p e c t i v e l y ,   a t   w h i c h   t h e   c o n f i g u r a t i o n a l   e n t r o p y   v a n i s h e s .   S i g n i f i c a n c e   o . f   l o w t e m p e r a t u r e   e l e c t r o l y t e   s o t u   t i o n   b e h a v i o r . A n   i m p o r t a n t   r e s u l t   o f   t h i s   t r e a t m e n t   w h i c h   w e   w i s h   t o   e m p h a s i z e   i n   t h i s   p a p e r   i s   t h e   r e a l i z a t i o n   t h a t ,   p r o v i d e d   c r y s t a l l i z a t i o n   h a s   n o t   a l   r e a d y   o c c u r r e d ,   a n y   e l e c t r o l y t e   s o l u t i o n   w i l l   l o s e   i t s   l i q u i d   c h a r a c t e r   a n d   b e c o m e   a   g l a s s   a t   a   ~ e m p e r a t u r e   ( ~ T o )   w h i c h   i n   g e n e r a l   w i l l   b e   w e l l   a b o v e   1 0 0 ~   W h e r e   t h e   e l e c t r o l y t e   p r o p e r t y   o f   i n t e r e s t   d e p e n d s   o n ,   o r   i s   r e l a t e d   t o ,   t h e   f l u i d i t y   o f   t h e   s u b s t a n c e ,   t h i s   t e m   p e r a t u r e   t h e r e f o r e   p l a c e s   a n   a b s o l u t e   l o w e r   l i m i t   t o   t h e   s e r v i c e a b i l i t y   o f   t h e   m a t e r i a l .   1   F u r t h e r m o r e ,   i t   i s   f o u n d   t h a t ,   f o r   m a n y   e l e c t r o l y t e   m i x t u r e s   a n d   s o l u t i o n s   a t   t e m p e r a t u r e s   n o t   t o o   f a r   a b o v e   T o ,   T o   i t s e l f   i s   t h e   o n l y   i m p o r t a n t   v a r i a b l e   i n   t h e   t r a n s p o r t   e q u a t i o n ,   s o   t h a t   t h e   l o w t e m p e r a t u r e   t r a n s p o r t   p r o p e r t i e s   o f   s u c h   l i q u i d s   a r e   t o   a   l a r g e   e x   t e n t   k n o w n   o n c e   T o   i s   k n o w n   ( 3 ) .   T h u s ,   t h e   s c a l i n g   f a c t o r   t o   b e   u s e d   i n   c o m p a r i n g   a   g i v e n   p r o p e r t y   a m o n g s t   d i f f e r e n t   s o l u t i o n s   i s ,   f r o m   E q .   [ 1 ] ,   ( T   T o ) .   T o   i l l u s t r a t e   t h e   u s e f u l n e s s   o f   t h i s   p o i n t ,   w e   t a k e   a n   e x a m p l e   f r o m   t h e   w o r k   o f   G a r r e t t   e t   a l .   ( 7 )   w h o   w e r e   i n v e s t i g a t i n g   v a r i o u s   s t r o n g   ( p r e s u m e d   f u l l y   d i s s o   c i a t e d )   e l e c t r o l y t e   s o l u t i o n s   f o r   s u i t a b i l i t y   i n   l o w   t e m p e r a t u r e   b a t t e r y   a p p l i c a t i o n s .   T h e i r   ( u n i n t e r   p r e t e d )   d a t a   o n   t h e   v i s c o s i t y   o f   s o m e   s o l u t i o n s ,   r e l a   t i v e   t o   t h e   v i s c o s i t y   o f   w a t e r   a t   2 5 ~   a r e   r e p r o d u c e d   i n   F i g .   l ( i ) .   T h e   s t r i k i n g   f e a t u r e   o f   t h e   d a t a   i s ,   o f   c o u r s e ,   t h e   v e r y   r a p i d   r i s e   i n   s o l u t i o n   v i s c o s i t y   a t   t h e   l o w e r   t e m p e r a t u r e s .   I n   F i g .   1   ( i i )   w e   s h o w   h o w   t h e   u s e   o f   a   T o   v a l u e   a p p r o p r i a t e   t o   e a c h   s o l u t i o n   i n   t h e   s c a l   i n g   f a c t o r   ( T   T o )   r e d u c e s   t h e   d a t a   a p p r o x i m a t e l y   t o   a   s i n g l e   c u r v e .   T h e   r e m a i n i n g   m i n o r   d i f f e r e n c e s   s e e m   t o   b e   d u e   m a i n l y   t o   v a r i a t i o n s   i n   t h e   v a l u e   o f   t h e   p r e e x p o n e n t i a l   t e r m   A ,   a s   s e e n   i n   t h e   f o l l o w i n g .   A c c o r d i n g   t o   E q .   [ 1 ] ,   t h e   v a l u e s   o f   T o   w h i c h   r e d u c e   t h e   s o l u t i o n   v i s c o s i t i e s   a s   i n   F i g .   l ( i i ) ,   s h o u l d   y i e l d   a   l i n e a r   p l o t   f o r   t h e   r e l a t i v e   v i s c o s i t i e s   w h e n   l o g   ~ / ~ ] o ,   ( o r   l o g   ~ o / ~   t o   m a i n t a i n   E q .   [ 1 ]   s i g n s )   i s   p l o t t e d   a g a i n s t   1 / ( T T o ) .   T h e   a p p r o p r i a t e   s e m i l o g a r i t h m i c   p l o t s   a r e   s h o w n   i n   F i g .   2 .   T h e   v a r i o u s   p l o t s   a r e   n o w   s e e n   t o   b e   d i f f e r e n t i a t e d   b y   s m a l l   c h a n g e s   i n   t h e   v a l u e s   o f   t h e   p a r a m e t e r   A ,   t h e   p l o t s   b e i n g   l i n e a r   w i t h   e s   s e n t i a l l y   e q u a l   v a l u e s   o f   t h e   s l o p e   k .   I t   m u s t   b e   s a i d ,   h o w e v e r ,   t h a t   t h e   a v a i l a b l e   d a t a   a r e   n o t   s u f f i c i e n t l y   T h e   l i m i t i n g   t e m p e r a t u r e   u s u a l l y   i m p o s e d   b y   t h e   c r y s t a l l i z a   t i o n   t e m p e r a t u r e   i s   n o t   a n   a b s o l u t e   l i m i t   i n s o f a r   a s   s u i t a b l e   a d d i   t i v e s   c a n   u s u a l l y   m a k e   c r y s t a l   n u c l e a t i o n   a   v e r y   i m p r o b a b l e   p r o c e s s   e v e n   w h e n   t h e r m o d y n a m i c a l l y   f a v o r e d .   \\x0c']"
},{
  "_id": 94,
  "PDF": "High‐Temperature Oxidation of ZrB 2 –SiC–AlN Composites at 1600°C.pdf",
  "Text": "['High-Temperature Oxidation of ZrB2  -SiC-AlN Composites at 1600°C  Gaoyuan Ouyang,  ‡,§  Pratik K Ray,  ‡,§  Matthew J Kramer,  ‡,§  and Mufit Akinc  ‡,§,†  ‡  Department of Material Science and Engineering, Iowa State University, Ames, Iowa  §  Division of Materials Science and Engineering, Ames Laboratory, Ames, Iowa  The effect of AlN substitution on oxidation of ZrB2-SiC was evaluated at 1600°C up to 5 h. Replacement of ZrB2 by AlN, with 30 vol% SiC resulted in improved oxidation resistance  with a thinner  scale and reduced oxygen affected area. On the  other hand, substitution of AlN for SiC resulted in a deteriora tion of the oxidation resistance with an abnormal scale and sig nificant  recession. The effect of SiC content was also studied,  and was found to be consistent with the literature for the com posites without AlN additions. A similar  effect was  observed  when AlN was added, with the  higher SiC content materials  showing  improved  oxidation  resistance. X-ray  photoelectron  spectroscopy showed the presence of Al2O3 and SiO2 on the surface, which could possibly lead to a modification in the vis cosity of the glassy oxide scale. Possibly, ior of ZrB2-SiC composites AlN additions by adjusting the Al:Si ratios.  the oxidation behav can be  improved with controlled  I.  Introduction  ZrB2-SiC composites high-temperature ceramics  are  promising  candidates  for  ultra for  hypersonic  applications  because of their unique thermal, chemical, and physical prop erties.  These composites have high melting temperatures, from 2270°C (ZrB2-SiC eutectic) to 3050°C1 ZrB2), low densities ranging from 3.21 g/cm3 6.08 g/cm3 60 W\\x01(m\\x01K) for ZrB2, and high thermal \\x001 at room temperature2). In order sify ZrB2-SiC composites at ambient pressure, tives such as WC, C, and/or B4C have been additions have also been  ranging  (pure  for  SiC to  conductivity  (c.a.  to fully den several addiproposed.3-8 considered and  Several  alloying  tested with the goal of ZrB2-SiC. Most microstructure and  improving the oxidation resistance of  of  these  studies  attempted  to modify  the  composition. The  common  approaches et al.9 as: increase the viscosity of the borosilicate glass (W,9 TaSi2 10); inhibit ZrO2 polymorphic transformation (Ta9); (3) substitute SiC with another silicon-containing compound 11); (Ta5Si3 introduce high-temperature protective 12); tory phase (LaB6 (5) decrease the porosity of ZrO2 by liquid phase sintering of ZrO2 scale (WC9). Recently, AlN has been proposed as a sintering aid for hot pressing of ZrB2-SiC composites.13,14 Addition of AlN leads to the partial removal of B2O3 from the surface of ZrB2 particles, which in turn help with the densification process. The reaction between AlN and B2O3 results in the formation of Al2O3 and BN, which possibly lower degree of grain growth compared to the unmodified samples.13,14 However, despite the availability of  toward this problem were  summarized by Eakins  (1)  (2)  (4)  refrac contributed to a  ZrB2 ture on the role of AlN as a sintering aid, not much has been  litera reported on how AlN affects the oxidation behavior of ZrB2-SiC composites. Al2O3, which forms when AlN is oxidized, is a good oxygen barrier at moderate temperatures  and  lowers  the  viscosity  of  the  borosilicate  scale,  thereby  improving its ability to flow into pores and form a continuous protective oxide layer.15 Taking these into account, AlN  substitution is expected to alter the ZrB2-SiC composites. The present study attempts to address the oxidation behavior of ZrB2-SiC composites with AlN substitutions at 1600°C.  oxidation  behavior  of  II.  Experimental Procedure  ZrB2 (Grade B, ~2 lm particle size, H. C. Starck, Karlsruhe, Germany), SiC (Grade UF-10, ~1 lm particle size, H. C. Starck), and AlN (Grade C, ~1 lm particle size, H. C. Star ck) were used as raw materials in this study. Four composisynthesized: ZrB2-30 vol%SiC (ZS73, ZrB2-30 vol%SiC-10 vol%AlN addition), (ZSA631, AlN substitution for ZrB2), ZrB2-20 vol%SiC-10 vol% AlN (ZSA721, AlN substitution for SiC), and ZrB2-20 vol% SiC (ZS82, no AlN addition). These composites were sintered 1600°C  tions were  no AlN  from dense  compact  and  subsequently  oxidized  at  and their oxidation behavior was studied.  Samples were sintered by following the procedure described by Zhang et al.5 Ceramic powders were wet milled  in a plastic  jar using WC as milling media using roller mill  (Cole-Parmer Lab mill 8000, Vernon Hills,  IL) and methyl  ethyl ketone  (MEK) as  solvent  for 24 h. The powders were (\\x191%-2%) of WC from milling media which is believed to be beneficial as a sintering aid.5 The mixture was milled for another 24 h after binder  contaminated by  a  small  amount  (QPAC40, polypropylene carbonate) addition. Following wet evaporated at 40°C under vacuum. milling, the solvent was The sample was ground and sieved through a 300-lm sieve.  Cylindrical  samples were prepared by pressing powders, first  uniaxially, followed by cold isostatic pressing at 310 MPa. The binder was burned out at 600°C for 1 h in flowing argon  atmosphere before sintering. Finally,  samples were placed in  a BN-coated graphite  crucible  and sintered in an electrical  resistance  furnace  equipped with graphite heating  elements  (3060-FP20, Thermal Technology  Inc.,  Santa Rosa, CA).  The sintering profile included two 1-h isothermal holds at 1250°C and 1450°C under vacuum to remove the surface oxides.5 Samples were then sintered at 2000°C in flowing helium  atmosphere for 2 h.  Sintered  densities were measured  using  the Archimedes  method according to the ASTM standard B962-13 and con verted to relative  theoretical densities. Theoretical densities  were estimated using the rule of mixtures under  the assump tion that  there was no change in composition during sinter ing. Microstructures  and phase  assemblages of  the  sintered  samples were  studied using a JEOL 5910Lv (Tokyo,  Japan)  scanning electron microscope (SEM) and a Philips PANalyti cal (Almelo, Netherlands) X-ray diffractometer (XRD). Com position  mapping  was  done  using  a  FEI  Quanta-250  (Hillsboro, OR) SEM equipped with Oxford Aztec  energy T. Parthasarathy—contributing editor  Manuscript No. 36608. Received March 23, 2015; approved October 19, 2015.  †  Author to whom correspondence should be addressed. e-mail: makinc@iastate.edu  808  J. Am. Ceram. Soc., 99 [3] 808-813 (2016)  DOI: 10.1111/jace.14039  © 2015 The American Ceramic Society  Journal  \\x0c', 'dispersive  X-ray  analysis  system.  Cross-sections  of  the  samples were polished before analysis. The samples were etched with molten NaOH/KOH (1:1 molar ratio) at ~200°C to reveal the ZrB2 grain boundaries for estimation of the grain sizes. Measurements were done from three representa tive areas of each sample, and the grain sizes estimated using  an Image Analysis program (ImageJ).  Samples were placed on ZrO2 crucibles and oxidation tests were carried out in a box furnace at 1600°C in ambient air for  up to 5 h. They were introduced directly to the preheated fur nace for isothermal oxidation and removed after desired time  intervals. Three samples were tested for each time period and  their  specific mass  changes were  averaged.  Samples were  weighed before and after  the  test. The oxidized surface was  studied using SEM, XRD, and XPS, after which the oxidized  cross-section microstructures were studied. For surface XPS, the measurement was done using PHITM Physical Electronics  5500 Multitechnique ESCA system (Physical Electronics Inc.,  Chanhassen, MN)  with  monochromatic  AlKa  radiation  (1486.6 eV). The peak positions were determined with refer ence to the adventitious carbon peak at 284.6 eV. The atomic  concentration was  calculated by using the  sensitivity factors  provided with the PHITM acquisition software. The thickness  of the oxide scale was determined by measuring the thickness at 10 areas or more, at 100 lm intervals, along the scale.  III.  Results and Discussion  Figure 1  shows  the microstructure of  the  sintered samples.  The  relative  densities  of  the  composites were  found to be  99.4%,  98.4%,  93.5%,  and  95.5% for  ZS73  [Fig.1(a)],  ZSA631 [Fig.1(b)], ZSA721 [Fig.1(c)], and ZS82 [Fig. 1(d)],  respectively. The  bright  phase  in  these micrographs  corre sponds  to ZrB2, while and/or AlN. SiC and AlN could not be differentiated in the  the dark phase  corresponds  to SiC  backscattered images due to almost  identical Z contrast. The  presence of AlN and SiC,  therefore, was confirmed using X ray  diffraction.  Presence  of  the  constituent  elements  and  phases was further confirmed with EDS elemental mapping in Figs. 2(a)-(d).  as  shown  Substitution  of AlN for ZrB2 (ZrB2 grain size 6.2 \\x06 0.3 lm in ZS73 and 3.9 \\x06 0.4 lm in [Figs. 1(b) and 2(b), ZSA631] leads to a finer microstructure ZSA631) which is in agreement with the literature.13,14  ZSA721 shows a larger mass gain during oxidation relative  to ZS82, ZS73, and ZSA631, as  shown in Fig. 3.  It appears  that ZS73 and ZSA631 exhibit similar mass gain and seem to  level off around 5 h, whereas ZSA721 shows higher mass gain  initially and continues to gain weight almost at a linear rate.  Similarly, the ZS82 sample, despite a lower initial mass change,  exhibits a significant mass gain with time. The oxidation of these composites is likely to involve multiple reactions.16 Possi ble chemical reactions are tabulated in Table I. Reactions (1),  (2), and (4) would result in a mass gain, while reactions (3) and  (5) would result  in mass loss. Reaction (3) also represents leading to a Si-depleted subsurface.16  active oxidation of SiC,  Figure 3  reflects  the net  effect of  these  reactions. Hence,  a  microstructural  study of  the oxidized cross-sections was car ried out to analyze the oxidation behavior of these samples.  Figure 4 shows  the 5-h oxidized cross-section microstruc tures of  these composites. ZS73, ZSA631, and ZSA721 show  a degree of  similarity  in their oxidation behavior, with the  scale  comprising  three  layers. While  the  thickness of  these  layers differed for each sample,  the general nature remained  the  same. The  top  layer  exhibiting  a  dark  contrast  corre sponds  to  the  silica  scale.  ZS73  [Fig. 4(a)]  and  ZSA631  [Fig. 4(b)]  exhibit a thinner  continuous oxide  scale  in com parison to ZSA721 [Fig. 4(c)]. ZSA721 has numerous pores  and  discontinuities  within  the  oxide  scale.  Additionally,  ZSA721  sample  also shows  the presence of ZrO2 surface, extending well into the oxide  channels  perpendicular  to the  subscale. Multiple ZrO2 clusters could also be seen throughout the microstructure. It has been shown in the literature  (a)  (b)  (c)  (d)  Fig. 1.  Backscattered images of ZrB2-SiC-AlN composites, after sintering. (a) ZS73, (b) ZSA631, (c) ZSA721, and (d) ZS82.  March 2016  Oxidation of ZrB2-SiC-AlN  809  \\x0c', 'that SiO2 affords improved oxidation resistance in comparin ZrB2-SiC systems.18,19 According to Opeka ison to ZrO2 et al.,19 parabolic rate constant for ZrO2 is much higher than SiO2. Thus, oxygen transport in ZrO2 channels is expected to be much higher than the silica scale. Although volume diffu sion of oxygen through ZrO2 grain is limited resulted from its poor electronic conductivity,20 interfacial diffusion of oxy gen through ZrO2 interface is three to four orders magnitude faster.21 Furthermore, the zirconia formed on the scale is not dense—pores  are  present  in  the  zirconia  scale,  and  rapid  penetration  of  oxygen  is  also  possible  through  pores  and  cracks  in  the ZrO2 channels and clusters  channel. Hence,  the  presence  of ZrO2 instead of an uninterrupted silica scale  exacerbates  the  oxidation  of  the  composite. The  oxidized  microstructures  of  ZS82  shows  the  similar  layered  microstructures,  but  followed  by  a  large  and  irregular  Si-depleted region, as shown in Fig. 4(d).  It  can be  seen from the  composition maps  in Fig. 5 that  the various  layers  in ZSA631, ZS73, and ZSA721 are rather  regular and relatively planar, whereas ZS82 shows significant  irregularities, which is in agreement with the results reported et al.22 The Si map shown in Fig. 5(h)  by Williams  shows  these  irregularities  clearly,  especially  in  case  of  the  Si depleted region. The oxidation behavior  in this composite is  therefore expected to be relatively stochastic.  The microstructures  show the presence of an intermediate  layer of oxide underneath the silica scale for ZS73, ZSA631,  and ZS721. The  elemental distributions mapped using EDS  (a)  (b)  (c)  (d)  Fig. 2.  Layered EDS elemental maps of ZrB2-SiC-AlN composites, after bright phase has high concentration in Zr, B; the gray phase has high concentration in Si, C;  sintering.  (a) ZS73,  (b) ZSA631,  (c) ZSA721, and (d) ZS82. The  the dark phase has high concentration in Al, N.  Fig. 3.  Mass  change  of  samples  on  oxidation  in  air  at  1600°C.  Each data point  represents average of  three samples with error bars  indicated for the range of mass change.  Table I.  Mass Change Associated with Various Oxidation Reactions17  Number  Reactions  Mass per mole of  Dm, g  1  ZrB2(s) + 5 2O2(g)? ZrO2(s) + B2O3(l) SiC(s) + 3 2O2(g) ? SiO2(l) +  ZrB2  80  2  CO(g) SiC(s) + O2(g) ? SiO(g) + CO(g) AlN(s) + 3 4 O2(g) ? 2Al2O3(s) + 1 2N2 (g)17 B2O3(l) ?B2O3(g)  SiC  20  3  SiC  \\x0040  4  1  AlN  10  5  B2O3  \\x0069.6  810  Journal of the American Ceramic Society—Ouyang et al.  Vol. 99, No. 3  \\x0c', 'indicate  the presence of Zr, Si, and O, while  the backscat tered images suggest the presence of two phases. Hence, prelayer comprised ZrO2 + SiO2, which would be in accordance with the results reported in the literature.16,23  sumably,  this  The  layer underneath this mixed oxide scale exhibits  signifi cant silicon depletion. This Si-depleted region forms the final  layer of  the oxygen affected area. The mechanism of elsewhere.16 While  silicon  depletion  has  been  discussed  all  these  three compositions show an intermediate ZrO2 + SiO2 of comparable thickness, differences in the top layer and the subscale region. ZSA631 top layer of ~30 \\x06 10 lm, while the thickest silica layer,  layer  of  they exhibit significant  [Fig. 4(b)] has  the thinnest  ~120 \\x06 10 lm. ZSA721 [Fig. 4(c)] ~50 \\x06 5 lm. The thickness of ZS73 has a a significant variation, but the  has  silica  top  layer  thickness  of  the silica scale in ZS82 shows  relatively uniform regions of  the scale are comparable to ZSA631, 35 \\x06 10 lm. The and ZSA631, with the latter having a marginally thinner sub(65 lm vis\\x12a-vis 50 lm). However, ZSA721 scale shows a very thick silicon-depleted subscale, ~150 lm. The ZS82 sam showing a thickness of  subscale  region  is  comparable  for ZS73  ple also exhibited the layered structure. The presence of SiO2 in the SiO2 + ZrO2 layer is quite low that the layer boundary is not easily recognizable to eye. In this case, there is a signif icant variation in the Si-depleted region ranging from 80 to 150 lm. Based on the microstructure and EDS analysis,  the  net oxygen affected region for ZS73, ZSA631, ZSA721, and 110-180 lm,  ZS82  were  160,  130,  300,  and  respectively  (shown schematically in Fig. 6).  X-ray data were  collected from the oxidized surfaces of  these samples (Fig. 7). The diffractograms from ZSA721 and  ZS82 were  dominated  by  the  peaks  of monoclinic ZrO2. 19° < 2h < 22°, angles  However,  at  the  lower  diffraction  extremely  small  peaks  corresponding  to  SiO2 or mullite were  can  be  observed.  Peaks  corresponding  to  zircon  absent. On the other hand, both ZS73 and ZSA631 exhibit high background in 2h = 20°-30° region which is the charac teristic of an amorphous phase. The nature of these diffraction patterns can be attributed to two factors—(i) Zr, having  a higher atomic number than Si (40 vis\\x12a-vis 14), has a higher  atomic scattering factor which translates intensities24 and (ii)  to stronger X-ray  the relative amount of ZrO2 at in ZSA721 or ZS82, compared to ZS73 and  the sur face  is higher  ZSA631. Crystalline ZrO2 will compared to SiO2, and this effect will be amount of ZrO2 increases indicated the presence of  result  in  a  higher  intensity  amplified as  the  relative to SiO2. The SEM studies significant amount of ZrO2 channels and clusters near the surface, embedded in the surround ing silica in ZSA721,  in comparison to ZS73 and ZSA631.  The SEM results are,  therefore,  further corroborated by the  X-ray evidence.  The  presence  of ZrO2 oxygen in ZSA721. This  channels would  provide  an  easy  pathway  for  results  in  higher  Si  recession in the subscale. However,  the presence of a signifi cant amount of ZrO2 is indicative of erage by the oxide scale in this particular  inadequate surface cov case.  Since  the  primary protection against oxidation in these materials is due  to the formation of continuous SiO2 layer, that the oxidation resistance would be significantly dependent  it stands to reason  on the amount of Si available to form SiO2. Furthermore, a lower SiO2 content in the scale is likely to be balanced by higher ZrO2 content, which has an adverse effect on the oxidation resistance since the presence of ZrO2 in the scale results in a discontinuous SiO2 layer. The effect of SiO2 content is also evident when the oxidation behavior of ZS73 and  ZS82 is compared. ZS82 has a larger  (and non-uniform) Si depleted region underneath the silica scale. This deterioration  in  oxidation  resistance caused by reduced literature.22 The  SiO2 degree  content  is  consistent with  the  lower  of mass  change in ZS82 is  likely due to the presence of higher ZrB2 content, which results in larger amounts of the volatile B2O3. lower net mass  This  contributes  to  larger mass  loss  and  gains. As discussed in the  text, mass  change  is not  the best  indication of oxidation resistance,  especially  since  it  repre sents  the  net  results  of  gain  (oxidations  remaining  on  the  scale) and loss  (volatile oxides). ZS73 and ZSA631 have the  same  volume  fraction  of  SiC. However, ZSA631  contains  10% AlN and 60% ZrB2 (as opposed to 70% ZrB2 in ZS73).  (a)  (c)  (d)  (b)  the 1600°C, 5-h oxidized coupons—(a) ZS73, Fig. 4. Cross-section microstructures of the structure of the oxide scale at a higher magnification (3009).  (b) ZSA631,  (c) ZSA721, and (d) ZS82. The inset shows  March 2016  Oxidation of ZrB2-SiC-AlN  811  \\x0c', '812  Journal of the American Ceramic Society—Ouyang et al.  Vol. 99, No. 3  (a)  (c)  (b)  (d)  (e)  (f)  (g)  (h)  Fig. 5.  EDS maps  of  corresponding  elements  for  four  5-h  oxidized  coupons.  (a) ZS73,  (b) ZSA631,  (c) ZSA721,  and  (d) ZS82; Their  corresponding Si maps are shown in e, f, g, and h.  (a)  (b)  (c)  (d)  Fig. 6.  Schematic  representation of  the oxygen affected area after  5 h oxidation, based on EDS mapping ((a) ZS73,  (b) ZSA631,  (c)  ZSA721, and (d) ZS82). The  light gray channels/cluster  in the  top  layer are ZrO2 channels.  that  scale  The oxide  initially forms at high temperatures scale. The SiO2-B2O3 these composites is a borosilicate system shows a low melting liquid around 440°C on the B2O3 rich side.25 As the SiO2 content increases, the liquidus temperature increases correspondingly. B2O3 is known to volatilize above 1200°C. Therefore, during the oxidation process, 1200°C, B2O3 evaporate from the surface, which drives the liquidus temper temperature  increases  above  starts  as  the  in  to  atures higher. This is likely to result  in increased viscosity of  Fig. 7. X-ray diffractogram of the surface of ZS73, ZSA631, ZSA721, and ZS82 coupons after oxidation in air at 1600°C for 5 h.  the oxide  scale, which in turn, will  result  in poorer  surface  coverage. Accurate quantification of boron is difficult. Mini mum amount of boron is expected in the scale surface based  on  results  the XPS  (Table II). This result with the results of Shugart et al.,26 Rezaie et al.,28 and Karlsdottir et al.29  that  is in agreement et al.,27 Levine  the surface B content  in  \\x0c', 'the glassy scale is low due boria,30 although its amount  to  the  high  vapor  pressure  of  could be higher deeper  in the  glassy layer.  Urbain  et al.  reported  the  viscosities  in  the Al2O3-SiO2 temperatures.15  system over  a  range  of  compositions  and  They showed that as the Al2O3 content increased, the viscosity of the Al2O3-SiO2 system decreased at 1600°C. The presof Al2O3 from the oxidation of AlN in ZSA631 is therefore likely  ence  to lower  the viscosity of  the oxide  scale  in  comparison to the ZS73 sample and hence result  in a contin uous  surface coverage. This,  in fact,  is  reflected in the cross section  micrographs  [Figs. 4(a)  and  (b)]  for  ZS73  and  ZSA631,  respectively, where  it  can be  clearly seen that  the  top layer  in oxidized ZS73 has a higher ZrO2 pared to ZSA631, which would have formed before complete  content  com scale  coverage was obtained. The presence of ZrO2 detected primarily because the depth probed by XPS is  is not  typi cally of  the order of  tens of nanometers, whereas  the ZrO2 depth of few  islands  in  the micrographs  show up  at  a  microns. As  expected,  given  the  similar AlN content,  the  XPS results  (Table  II)  from the oxidized ZSA721 showed a  similar  surface Al  content. However, a greater volume  frac tion of ZrB2 leads to greater availability of Zr to form ZrO2, in comparison to ZSA631. The porous and permeable ZrO2, once formed in significantly larger quantities, compared to  ZSA631, provides oxygen pathway deeper  into the material,  which  accounts  for  the  difference  in  oxidation  resistance  between these two composites.  IV.  Conclusions  ZrB2-SiC-AlN compacts with and without AlN were synthesized by pressureless sintering. AlN additions resulted in finer  microstructures,  when  substituted  for  ZrB2. affected region showed a degree of similarity with a SiO2-rich layer forming on the surface, followed by a ZrO2 + SiO2 interlayer, and finally a Si-depleted subscale. Chemical analy The  oxygen  sis  showed the presence of Al  (possibly as Al2O3) surface. Based on the work  in minor  quantities on the Urbain et al.,15  oxidized  of  this  is  likely  to reduce  the  viscosity of  the  protective SiO2-rich layer, which can flow and cover the surface, thereby protecting the composite from further oxidation. The ZrB2-30 vol%SiC-10 vol%AlN (ZSA631) showed the best oxidation resistance while the ZrB2-20 vol%SiC- the worst oxidation resis 10 vol%AlN (ZSA721)  showed  tance  due  to  the  copious  formation  of  oxygen  permeable  ZrO2 tion,  channels  and clusters,  excessive  silica  viscosity  reduc limited Si supply, and higher initial porosity. Therefore,  we conclude that a more oxidation resistant material can be  made with a proper Al  to Si  ratio,  tailoring the glass phase  viscosity while still providing sufficient Si content  to form a  protective oxide scale.  Acknowledgments  This work was supported by the AFOSRHTAM under contract # FA9550-11 1-201. The  authors  greatly  acknowledge  the  support  from this project. The  authors thank Eric Neuman and Prof. William G. Fahrenholtz for their assistance regarding the pressureless sintering of ZrB2-SiC. The authors also thank Warren Straszheim for his help in the EDS mapping analyses and Jim Ander egg for the XPS analyses.  References  1A. McHale, H. McMurdie, and H. Ondik, Phase Equilibria Diagrams. Vol ume X, Borides, Carbides, and Nitrides  : Phase Diagrams  for Ceramists. Vol.  X. American Ceramic Society, Westerville, Ohio, 1994. 2J. W. Zimmermann, G. E. Hilmas, W. G. Fahrenholtz, R. B. Dinwiddie, W. D. Porter, and H. Wang, “Thermophysical Properties of ZrB2 and ZrB2- SiC Ceramics,” J. Am. Ceram. Soc., 91 [5] 1405-11 (2008). 3A. Chamberlain, W. Fahrenholtz, and G. Hilmas, “Pressureless Sintering of Zirconium Diboride,” J. Am. Ceram. Soc., 89 [2] 450-6 (2006). 4W. Fahrenholtz, G. Hilmas, S. Zhang, and S. Zhu, “Pressureless Sintering  of Zirconium Diboride: Particle Size Soc., 91 [5] 1398-404 (2008). 5S. C. Zhang, G. E. Hilmas, and W. G. 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Lorincz, “Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,” J. Mater. Sci., 39 [19] 5969-77 (2004). 11I. Talmy, J. Zaykoski, M. Opeka, and A. Smith, “Properties of Ceramics in the System ZrB2-Ta5Si3,” J. Mater. Res., 21 [10] 2593-9 (2006). 12X. Zhang, P. Hu, J. Han, L. Xu, and S. Meng, “The Addition of Lan a-Struc.  thanum Hexaboride to Zirconium Diboride tance,” Scripta Mater., 57 [11] 1036-9 (2007). 13F. Monteverde and A. Bellosi, “Beneficial Effects of AlN as Sintering Aid  for  Improved Oxidation Resis on Microstructure and Mechanical Properties of Hot-Pressed ZrB2,” Adv. Eng. Mater., 5 [7] 508-12 (2003). 14W. Han, G. Li, X. Zhang, and J. Han, “Effect of AlN as Sintering Aid on Hot-Pressed ZrB2-SiC Ceramic Composite,” J. Alloy. Compd., 471[1-2], 488- 91 (2009). 15G. Urbain, Y. Bottinga, and P. Richet, “Viscosity of Liquid Silica, Silicates and Alumino-Silicates,” Geochim Cosmochim Ac, 46 [6] 1061-72 (1982). 16W. G. Fahrenholtz, of ZrB2-SiC Oxidation: 143-8 J. Am. Ceram. Soc., 90 [1]  “Thermodynamic Analysis  Formation  of  a  SiC-Depleted Region,”  (2007). 17P. Boch, J. C. Glandus, J. Jarrige, J. P. Lecompte, and J. Mexmain, “Sin tering, Oxidation and Mechanical Properties Nitride,” Ceram. Int., 8 [1] 34-40 (1982). 18W. C. Tripp, H. H. Davis, and H. C. Graham, “Effect of an SiC Addition on Oxidation of ZrB2,” Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 19M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, “Oxidation-Based Materials Selection for 2000°C Plus Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience,” J. Mater. Sci., 39 [19] 5887-904 (2004). 20T. A. Parthasarathy, R. A. Rapp, M. Opeka, and R. J. KeranS, “A Model the Oxidation of ZrB2, HfB2 and TiB2,” Acta Mater., 55 [17] 5999-6010 (2007). 21U. Brossmann, R. Wurschum, U. Sodervall, and H. E. Schaefer, “Oxygen  of Hot  Pressed Aluminium  for  Diffusion in Ultrafine Grained Monoclinic ZrO2,” 7646-54 (1999). 22P. A. Williams, R. Sakidja, J. H. Perepezko, and P. Ritt, “Oxidation of ZrB2-SiC Ultra-High Temperature Composites over Content,” J. Eur. Ceram. Soc., 32 [14] 3875-83 (2012). 23A. Rezaie, W. Fahrenholtz, and G. Hilmas, “Evolution of Structure Dur J. Appl. Phys.,  85  [11]  a Wide Range  of  SiC  ing the Oxidation of Zirconium Diboride-Silicon 1500°C,” J. Eur. Ceram. Soc., 27 [6] 2495-501 (2007). 24B. Cullity  Carbide  in Air  up  to  and  S.  Stock, Elements  of X-Ray Diffraction.  Prentice-Hall,  Upper Saddle River, New Jersey, 2001. 25T. J. Rockett and W. R. Foster, “Phase Relations Oxide-Silica,” J. Am. Ceram. Soc., 48 [2] 75-80 (1965). 26K.  in the System Boron  Shugart,  S. Y. Liu, F. Craven, and E. Opila, “Determination of Retained B2O3 Content in ZrB2-30 vol% SiC Oxide Scales,” J. Am. Ceram. Soc., 98 [1] 287-95 (2015). 27A. Rezaie, W. Fahrenholtz, and G. Hilmas, “Oxidation of Zirconium Diboride-Silicon Carbide at 1500°C at a Low Partial Pressure of Oxygen,” J. Am. Ceram. Soc., 89 [10] 3240-5 (2006). 28S. Levine, E. Opila, M. Halbig, J. Kiser, M. Singh, and J. Salem, “Evalua tion of Ultra-High Temperature Ceramics Ceram. Soc., 22 [14-15] 2757-67 (2002). 29S. N. Karlsdottir and J. W. Halloran, “Rapid Oxidation Characterization 3233-8  for Aeropropulsion use,”  J. Eur.  of Ultra-High Temperature Ceramics,”  J. Am. Ceram. Soc.,  90  [10]  (2007). 30R. H. Lamoreaux, D. L. Hildenbrand, and L. Brewer, “High-Temperature  Vaporization Behavior of Oxides .2. Oxides of Be, Mg, Ca, Sr, Ba, B, Al, Ga, in, Tl, Si, Ge, Sn, Pb, Zn, Cd, and Hg,” J. Phys. Chem. Ref. Data, 16 [3] 419-  43 (1987).  h  Table II.  XPS Data from the Oxidized Surface (Units in  Atomic Percentage)  Sample  B  O  Al  Si  Zr  Al:Si  B:Si  ZS73  2  71  —  26  <1 <1 <1  —  0.1  ZSA631  2  72  4  22  0.2  0.1  ZSA721  3  72  4  20  0.2  0.15  ZS82  2  71  —  23  5  —  0.1  March 2016  Oxidation of ZrB2-SiC-AlN  813  \\x0c']"
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  "_id": 95,
  "PDF": "Hot pressing and oxidation behavior of ZrB2–SiC–TaC composites.pdf",
  "Text": "['Ceramics International 46 (2020) 3725-3730  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Hot pressing and oxidation behavior of ZrB2-SiC-TaC composites  Mahdi Ghassemi Kakroudia,∗, Masoumeh Dehghanzadeh Alvaria, Mehdi Shahedi Aslb, Nasser Pourmohammadie Vafaa, Taher Rabizadeha  T  a Department of Materials Science and Engineering, University of Tabriz, Tabriz, Iran b Department of Mechanical Engineering, University of Mohaghegh Ardabili, Ardabil,  Iran  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ultrahigh temperature ceramics Microstructure Tantalum carbide Oxidation rate Sintering  A ZrB2-20 vol% SiC-5 vol% TaC ceramic was sintered through the hot pressing process. By sintering for 60 min at 1850 °C under 40 MPa in a vacuum atmosphere, an almost fully dense specimen with a relative density of 97.5% was achieved. The phase and microstructural investigations revealed that alongside the ZrB2, SiC and TaC as the starting materials, two new phases of ZrC and TaSi2 were synthesized in-situ during the hot pressing. The oxidation behavior of as-sintered hybrid composites was investigated at three temperatures in an atmospheric furnace for 1-, 4and 10-h cycles. Parabolic oxidation rate constants of 2.25, 30.69 and 1309.70 mg2/ cm4.h were calculated for the samples oxidized at the temperatures of 1000, 1400 and 1700 °C, respectively. The activation energy of 178.7 kJ/mol was kinetically calculated for the oxidation of ZrB2-SiC-TaC ceramics. Four oxidized layers on the unaffected ZrB2-SiC-TaC matrix were also detected and microstructurally characterized.  1.  Introduction  Ultrahigh temperature ceramics and composites (UHTCs) are nominated as promising candidates for aerospace applications including the nose-caps and leading-edges of the advanced hypersonic vehicles [1-13]. ZrB2-SiC-based ceramics, due to their high strength, high melting temperatures, high thermal conductivity, and relatively low density have recently attracted the consideration of researchers in this field. Such UHTCs usually serve at ultrahigh temperature environments; hence, the investigation on the oxidation behavior of them has also attracted considerable attentions [14-21]. Because of the formation of a glassy borosilicate protective layer, it is well-known that the SiC reinforced ZrB2-based composites have a higher oxidation resistance than that of the monolithic ZrB2 [22-33]. The oxidation manner of ZrB2-30 vol% SiC composite was studied under reducing situations at 1500 °C, and it was found that an oxide non-preserve surface scale is produced during the oxidation process in CO-CO2 system [34-41]. A multi-layer oxide scale, with a SiO2-rich borosilicate liquid, was formed during the oxidation of ZrB2-SiC UHTCs at the elevated temperatures (1550 and 1700 °C) in the ambient air. The deposition mechanism and nature of the ZrO2 was also investigated in the ZrB2-15 vol% SiC ceramic [42,43]. In another study, the oxide scale of the ZrB2-based materials, reinforced with 10 and 30 vol% SiC, was found to be composed of a glassy layer with an interior SiC-depleted zone [29,44-49]. A thermodynamic exemplar was also broadened for  the explanation of the appearance of such SiC-depleted zone throughout the oxidation of ZrB2-SiC materials at 1500 °C in the air, which suggested a structure consisting of SiO2-rich, Zr-plentiful oxides, as well as SiC-depleted layers [50]. By introducing the tantalum compounds (e.g. TaB2, TaSi2, and TaC), the oxidation resistance of ZrB2-SiC materials may be influenced, but several irreconcilable results have been reported. Adding TaSi2 led to an improvement in the oxidation resistance of ZrB2-SiC ceramics at < 1400 °C; however, the presence of tantalum disilicide worsened the oxidation behavior at higher temperatures [51]. The oxidation behavior of ZrB2-SiC composites was also enhanced through the introduction of TaSi2 and TaB2 additives at the temperature range of 1150-1550 °C. Anyway, detrimental influence on the oxidation resistance of ZrB2‐based UHTCs was observed at 1800 °C by the addition of Tiand Ta-based compounds [52]. It was also stated that the oxidation resistance of ZrB2-SiC UHTCs, between 1200 and 1500 °C, was reduced by the addition of 10 vol% TaC while incorporating 30 vol% TaC resulted in an intensified resistance against the oxidation [53]. In this work, a ZrB2-20 vol% SiC sample containing 5 vol% TaC additive is hot pressed for 60 min at 1850 °C under 40 MPa and its microstructure, densification and oxidation behavior at the temperatures of 1000, 1400 and 1700 °C an atmospheric furnace are investigated.  ∗ Corresponding author. E-mail address: mg_kakroudi@tabrizu.ac.ir (M. Ghassemi Kakroudi).  https://doi.org/10.1016/j.ceramint.2019.10.093 Received 10 September 2019; Received in revised form 10 October 2019; Accepted 10 October 2019  Available online 11 October 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'M. Ghassemi Kakroudi, et al.  Ceramics International 46 (2020) 3725-3730  Fig. 1. XRD pattern of the mixture of ZrB2/SiC/TaC powders.  Fig. 2. FESEM fractograph of  the hot pressed ZrB2-SiC-TaC composite.  Fig. 3. XRD pattern of the hot pressed ZrB2-SiC-TaC composite.  Fig. 4. Square of the specific weight change versus the oxidation time at (a) 1000 °C, (b) 1400 °C, (c) 1700 °C, and (d) plot of the logarithmic parabolic rate constants as a function of the reciprocal temperatures.  3726  \\x0c', 'M. Ghassemi Kakroudi, et al.  Ceramics International 46 (2020) 3725-3730  TaC as the starting materials, some peaks of ZrO2 are also detected by the phase analysis. The identification of ZrO2 is attributed to the oxide impurity which has formed on the surface of ZrB2 powders as the main component of the mixture. The formation of SiO2 and Ta2O5 on the surfaces of SiC and TaC particles, respectively, is also possible; however, due to the lower amounts of such oxides, they could not be detected by XRD test. the ZrB2-SiC-TaC A relative density of 97.5% was obtained for composite after hot pressing for 60 min at 1850 °C under 40 MPa. Although the ceramic composite has approached its theoretical density, the remaining 2.5% porosity in the as-sintered specimen can be related to insufficient sintering temperature or the absence of any metallic, carbonaceous or other types of sintering aids. Anyway, due to the negative influences of employing the sinter additives or higher sintering temperatures on the mechanical performance of sintered parts, a value of 97.5% seems to be acceptable. The FESEM image of the fracture surface of the as-hot-pressed composite is shown in Fig. 2. The grey, dark, and bright areas in this fractograph are ZrB2, SiC, and TaC grains, respectively, which is somehow consistent with the volume ratios of each phase in the powder mixture. Moreover, such a microstructural outcome is in good harmony with the relative density value of 97.5% for the sintered bulk. the hot pressed ZrB2-SiC-TaC Fig. 3 exhibits the XRD pattern of composite. Although the presence of some oxide impurities in the asmixed powders is previously identified by XRD (Fig. 1), no noticeable oxides can be found in the X-ray diffraction pattern of the as-sintered sample (Fig. 3). The XRD analysis verified the in-situ formation of two new compounds after the hot pressing process, namely the ZrC and TaSi2 phases. A hypothetical chemical reaction suggested as Eq. (1), may justify the results of XRD test. The progression of such reaction (Eq. (1)) not only leads to the formation of TaSi2 and ZrC phases but also results in the disappearance of ZrO2 as the surface impurity. The investigation of this reaction by the HSC Chemistry package disclosed that Eq. (1) can be favorable at the temperatures above 2078 °C, under the standard atmospheric state (1 atm). Thermodynamically, this reaction cannot occur at the hot pressing temperature of 1850 °C at first glance. However, it should be noted that the sintering process was conducted under the vacuum level of 5 × 10−2 Pa. Hence, regarding the mentioned fact and due to the formation of a gaseous byproduct (carbon monoxide), the suggested reaction of Eq. (1) seems to be progressed at lower temperatures.  2SiC + TaC + ZrO2 = TaSi2 + ZrC + 2CO (g)  (1)  Wang et al. [53] recently showed that the oxidation of ZrB2-SiC-TaC composites are controlled by the diffusion process; hence, the oxidation obeys a parabolic behavior according to Eq. (2):  w2 = k.t  (2)  where w is the specific weight change, k is the parabolic oxidation rate constant, and t is the oxidation time. The plots of the square of the specific weight variations versus the oxidation time at the temperatures of 1000, 1400 and 1700 °C are shown in Fig. 4a-c, respectively. The parabolic oxidation rate constants, at different oxidation temperatures, are calculated from the slopes of the trend-lines: k 1000 °C = 2.25 mg2/ cm4.h, k 1400 °C = 30.69 mg2/cm4.h, and k 1700 °C = 1309.70 mg2/cm4.h. The activation energy of the oxidation process can be estimated from the parabolic oxidation rate constants assessed at the various temperatures employing the Arrhenius formula (Eq. (3)):  k\\xa0  =  \\xa0A\\xa0exp\\xa0  ΔE RT  ⎛ ⎝  ⎞ ⎠  (3)  Fig. 5. XRD pattern of  the oxidized ZrB2-SiC-TaC UHTC at 1700 °C.  2. Experimental method  2.1. Materials and method  A ZrB2-based UHTC, with 20 vol% SiC and 5 vol% TaC additives, was fabricated by hot pressing with no sintering aid. The purchased raw size < 2 μm, materials were ZrB2 powder (purity > 99%, Chinese Leung Hi-tech Co.), SiC powder (purity > 99%, size < 0.7 μm, Chinese size < 0.5 μm, Leung Hi-tech Co.) and TaC powder (purity > 99%, Chinese Huhan Haiyun Metallurgical Materials Co.). At first, these powders with the designated volume ratios were mixed for 2 h in a polyethylene cup by ZrO2 balls with alcohol (96%) as the mixing medium. After the ball-mixing process, the powder mixture was dried in an oven at 120 °C for 12 h. The resultant dehumidified mixture was then put into a graphite mold, covered by flexible graphite foils, and hot pressed (Nanozint 10i, Khala Poushan Felez, Iran) at 1850 °C for 60 min under a uniaxial load of 40 MPa in 5 × 10−2 Pa vacuum with a heating rate of 10 °C/min. The oxidation investigations were performed isothermally at 1000, 1400 and 1700 °C for 1-, 4and 10-h cycles in an atmospheric furnace (Azar Furnaces. Co., Iran). Three specimens used for the oxidation test were prepared by wire-cut electrical discharge machining in the dimensions of 15 × 5 × 5 mm3.  2.2. Characterization  The rule of mixtures and the Archimedes methodology were utilized for the estimation of theoretical, bulk and relative density values of hot pressed UHTC composite. The as-mixed powders, the as-sintered composite and the as-oxidized sample at 1700 °C were phase-analyzed via an X-ray diffractometer (German Siemens D5000). The microstructural characterizations were carried out employing a field-emission scanning electron microscope (Tescan Mira3, Czech Republic) which was equipped with and energy dispersive spectroscope (Digital X-ray processor: DXP-X10P). The weight gain at three oxidation temperatures, as a function of the time of oxidation, was measured using a balance with 0.001 g accuracy (AND GE-220). The thermodynamic assessments of the chemical reactions, during the sintering and oxidation processes, were done using the HSC Chemistry package (version 5.11).  3. Results and discussion  Fig. 1 displays the XRD pattern of as-mixed ZrB2/SiC/TaC powders. As seen in this spectrum, beside the diffraction peaks of ZrB2, SiC and  where A is the constants, R is the universal gas constant, ΔE is the apparent activation energy, and T is the oxidation temperature in Kelvin. Fig. 4d plots the ln (k) as a function of the reciprocal temperatures; hence, the activation energy of 178.7 kJ/mol is calculated for the  3727  \\x0c', 'M. Ghassemi Kakroudi, et al.  Ceramics International 46 (2020) 3725-3730  Fig. 6. FESEM image and EDS spectra of the different layers in the cross section of oxidized ZrB2-SiC-TaC UHTC at 1400 °C.  3728  \\x0c', 'M. Ghassemi Kakroudi, et al.  Ceramics International 46 (2020) 3725-3730  oxidation of ZrB2-SiC-TaC composite in this research work. The oxidized surface of ZrB2-SiC-TaC ceramic at 1700 °C is investigated by XRD analysis and the results are shown in Fig. 5. It can be detected that the oxide layer contains the remarkable ZrO2 phase with the Ta2O5, SiO2, and ZrSiO4 phases. Such an outcome clearly verifies that all the components of the as-sintered composite including the starting particles of ZrB2, SiC, and TaC as well as the in-situ synthesized phases of ZrC and TaSi2 have been oxidized during the oxidation process. However, the complementary studies were carried out using the FESEM/EDS analyses in the following paragraphs. Fig. 6 displays the cross section of the ZrB2-SiC-TaC composite oxidized at 1400 °C temperature. It can be seen that there are four oxygen-affected layers on the unaffected ZrB2-SiC-TaC bed. The outer layer is a thin (< 10 μm) SiO2-rich layer containing the Ta2O5 phase. The adjoining inner layer includes the ZrO2, ZrSiO4, and SiO2 as the major phases with a thickness of ~20 μm. The third layer from the top surface of the oxidized sample seems to be a ZrO2-rich and SiC-depleted zone with an almost thickness of ~10 μm. The fourth layer, just above the unaffected ZrB2-SiC-TaC matrix, is a ~15-μm ZrB2-depleted layer contains SiO2 and Ta2O5 phases.  4. Conclusions  In conclusion, the hot pressing, microstructure and oxidation behavior of ZrB2-20 vol% SiC UHTCs with 5 vol% TaC were scrutinized. The sintering process was conducted for 60 min at 1850 °C under 40 MPa and the oxidation phenomenon was studied at 1000, 1400 and 1700 °C for 1, 4 and 10 h in the air. A near fully dense UHTC was manufactured through the hot pressing which the final microstructure contained the in-situ formed ZrC and TaSi2 phases beside the ZrB2, SiC and TaC as the initial components. The parabolic oxidation rate constants at different oxidation temperatures were estimated kinetically for ZrB2-SiC-TaC UHTCs, the oxidation of and activation energy of 178.7 kJ/mol was calculated. The microstructural investigations by SEM/EDS verified the formation of four distinctive oxide layers on the unaffected ZrB2-SiC-TaC matrix.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  References  [1]  [4]  S. Nekahi, F. Sadegh Moghanlou, M. Vajdi, Z. Ahmadi, A. Motallebzadeh, M. Shahedi Asl, Microstructural, thermal and mechanical characterization of TiB2-SiC composites doped with short carbon fibers, Int. J. Refract. Metals Hard Mater. 82 (2019) 129-135, https://doi.org/10.1016/j.ijrmhm.2019.04.005. [2] M. Sakkaki, F. Sadegh Moghanlou, M. Vajdi, F. Pishgar, M. Shokouhimehr, M. Shahedi Asl, The effect of thermal contact resistance on the temperature distribution in a WC made cutting tool, Ceram. Int. (2019), https://doi.org/10.1016/j. ceramint.2019.07.241. [3] M. Dashti Germi, Z. Hamidzadeh Mahaseni, Z. Ahmadi, M. 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  "_id": 96,
  "PDF": "Improved aero-thermal resistance capabilities of ZrB2-based ceramics in hypersonic environment for increasing SiC content.pdf",
  "Text": "['Contents lists available at ScienceDirect   Corrosion Science   journal homepage: www.elsevier.com/locate/corsci   Improved aero-thermal resistance capabilities of ZrB2-based ceramics in  hypersonic environment for increasing SiC content   Stefano Mungiguerra a, *, Anselmo Cecere a, Raffaele Savino a, b, Federico Saraga b,  Fr´ed´eric Monteverde b, Diletta Sciti b   a University of Naples “Federico II”, Department of Industrial Engineering, Aerospace Division, P.le Tecchio 80, 80125 Napoli, Italy  b National Research Council of Italy, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy     A R T I C L E I N F O   Keywords:  A. Ceramic  A. Zirconium  B. SEM  C. Oxidation  C. High temperature corrosion  C. Thermodynamic diagrams   1.  Introduction   A B S T R A C T   The  effects  of  SiC  content  on  aero-thermal  resistance  capabilities  of  ZrB2-xSiC-5Y2O3  ceramics  (x = 5,10,15,18.5 vol%) in atmospheric re-entry environment were investigated, exposing four hot-pressed  button-like specimens to an arc-heated plasma wind tunnel supersonic flow (enthalpy up to 20 MJ/kg). As x  increased, the maximum steady-state surface temperature decreased. A 400 K-spontaneous temperature jump  was observed on the ZrB2-5SiC-5Y2O3 front surface, with final temperature exceeding 2750 K. The oxide scale  consisted in Y-stabilized tetragonal/cubic zirconia, borosilicate glass, and Y2Si2O7 in minor amount. SEM-EDS  documented the microstructural differences for varying x, correlating the aero-thermal behavior, including the  temperature jump, to borosilicate glass durability.     The  extremely  demanding  aero-thermo-dynamic  conditions  encountered by hypersonic vehicles during atmospheric re-entry make  the design of proper Thermal Protection Systems (TPS) a topic of utmost  importance, requiring continuous improvement of materials and technologies [1,2]. The re-entry environment includes hypersonic Mach  numbers, projected  service  temperatures well above 2300 K and  recombination  reactions of dissociated gases which  substantially  enhance the heat flux on the most exposed surfaces of the spacecraft [3,  4].  Ultra-high temperature ceramics (UHTCs) and, more recently, ultra high temperature ceramic matrix composites (UHTCMCs), based on IV-V  transition metals carbides and diborides, are actively studied as candidates for these applications due to melting temperatures amongst the  highest known, excellent retained strength [5] and very good oxidation  resistance along these temperature ranges where state-of-the-art materials definitely fail to survive [6,7].  Silicon carbide is often added as minor component into UHTC  matrices to confer improved performances [8]. The dispersion of SiC in  the form of particle, short fiber or whisker into the main UHTC matrix is  frequently used to improve mechanical strength and damage tolerance  [5]. Oxidation resistance is also enhanced thanks to the formation of a   protective oxide scale, performing a self-healing function at ultra-high  temperatures [9,10].  SiC-containing MeB2 ceramics or Cf-SiC based CMC taken as baseline,  respectively, of the UHTCs and UHTCMCs class of materials, have  demonstrated excellent capability to endure the aero-heating as long as  the surface temperatures do not exceed 2200 K [10,11]. However,  increasing the heat load (and therefore the surface temperature), a  sudden temperature overshoot (better known as spontaneous temperature jump) of up to 400 K occurring at constant specific total enthalpy  was repeatedly observed [12-19]. A number of interpretations were  proposed: the loss of protective silica glass which substantially increases the chemical component of heat flux delivered to the surface  [12]; a transition from passive to active oxidation of SiC [13-15]; an  enhanced catalytic recombination of N2 due to the presence of Si(g)  [16]; an accelerated oxygen diffusion rate through new formed cracks  which promotes exothermic reactions of oxidation and nitridation [15,  17]; a surface chemical modification which alter emissivity and catalycity [18,19].  It has also been demonstrated [20] that, for increasing total enthalpies, the amount of dissociated oxygen and nitrogen increases,  making the catalytic effect more relevant in the determination of the  heat fluxes, which could be one of the reasons for the dependence of the  spontaneous temperature jump on the flow total enthalpy. Similar   * Corresponding author.  E-mail address: stefano.mungiguerra@unina.it (S. Mungiguerra).    https://doi.org/10.1016/j.corsci.2020.109067  Received 11 June 2020; Received in revised form 5 October 2020; Accepted 8 October 2020     CorrosionScience178(2021)109067Availableonline14October20200010-938X/©2020ElsevierLtd.Allrightsreserved.    \\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             features had moreover already been observed on fiber-reinforced UHTCs  tested in the same conditions [21]. Based on numerical simulations, the  temperature jump was attributed to a twofold mechanism, related to the  appearance of a zirconia scale upon the front surface: an increase in the  chemical contribution to the heat flux, due to an increase of surface  catalycity, and a reduction of the thermal conductivity through the  exposed oxide layer. The decrease in the thermal conductivity could  justify the concentration of the ultra-high temperatures in the front  surface, with a substantial heat transfer decay to the rear part of the  material.  Despite the variety of interpretation, a sound consensus on one of  them has not been reached yet. In this perspective, the first feature of  novelty of the present work was to carry out a focused testing campaign  to shed new light on the aero-thermal resistance capabilities of SiC containing ZrB2-based ceramics, with a specific interest on the effect  of SiC concentration. Lab-scale specimens with different SiC amount  (5-18.5 vol. %) were tested in high-enthalpy arc-jet wind tunnel conditions, where Earth’s entry conditions were reproduced. An important  element of novelty lied in the systematic approach to use the same setup  in terms of specimen geometry and power ramping, letting the only ZrB2  to SiC volume ratio vary. Non-intrusive diagnostic tools were jointly  employed to monitor the surface temperature of the specimen. The effects of increasing SiC contents on the material behavior in terms of  oxidation resistance, maximum equilibrium temperature, surface spectral emissivity and other unexpected phenomena such as spontaneous  temperature jumps were addressed. Another feature of novelty was the  use of a phase like Y2O3 to induce the stabilization of the zirconia and  therefore prevent the tetragonal/cubic to monoclinic phase transition  upon fast cooling. The same amount of 5 vol.% was added to various  ZrB2-SiC compositions: its effects were also investigated.   2. Materials and method   2.1. Materials and characterizations   Four ultra-high temperature ceramics, all based on a ZrB2 matrix,  were manufactured varying the ZrB2 to SiC volume ratio. The compositions (amount in vol. %) are labelled as    ZSY-05: ZrB2 + 5 SiC + 5 Y2O3                                                               ZSY-10: ZrB2 + 10 SiC + 5 Y2O3                                                              ZSY-15: ZrB2 + 15 SiC + 5 Y2O3                                                              ZSY-18: ZrB2 + 18.5 SiC + 5 Y2O3                                                          and were prepared by mixing commercial powders of ZrB2 (grade B,  H.C. Starck - Germany), SiC (UF25, H.C. Starck - Germany) and Y2O3  (grade C, H.C. Starck - Germany) using ethyl alcohol and SiC balls  mixing media. The mixtures were then dried using a rotary evaporator,  sieved (150 μm mesh size) and hot-pressed. Hot pressing was performed  in vacuum (10(cid:0) 100 Pa) heating up to 2173 K, 10 min isothermal dwell  at 2173 K, applying 30 MPa for the entire duration of the thermal  treatment.  The hot-pressed billets had a diameter of 32 mm and a thickness of  about 10 mm: they were ground on both sides to remove the residues of  graphite foil and reaction layers. The bulk density was measured by  Archimedes method using distilled water as the immersing medium.  From each hot-pressed billet a representative piece was cut and then  polished to 1 μm finish using diamond suspensions; the last step was a  finishing treatment using a colloidal silica suspension. Phase composition of the as-sintered (polished) material was determined by means of a  laboratory x-ray diffractometer (XRD, mod. D8 Advance, Bruker Germany), using a Ni filtered Cu Kα radiation. The microstructures were  also analyzed using a field emission scanning electron microscopy  (FESEM, ZEISS Sigma Germany) with a combination of secondary (SE)   and back-scattered electron (BSE) detector, and were integrated using  the energy dispersive spectrometer (EDS, INCA Energy 300, Oxford Instruments UK), equipped on the same FESEM.   2.2. Aero-heating set-up   To characterize the capabilities of various ZrB2 - x SiC - 5 Y2O3  compositions in an extreme aero-thermo-chemical environment, lab scale specimens were shaped as flat round button (Fig. 1a,b). Electro discharge machining was first used to extract the gross shape of the  specimen (pin included); then dimensions and roughness were finished  using conventional machining: the final roughness (Ra) of the front face  was about 0.2 μm.  The button-like specimens were exposed to a high-enthalpy supersonic dissociated air flow, typical of re-entry conditions, using the arc-jet  wind tunnel available at the University of Naples “Federico II”, named  SPES (Small Planetary Entry Simulator). This  is an open circuit,  continuous wind tunnel where a nitrogen plasma can be generated by an  industrial torch able to operate at powers up to 80 kW and mixed to a  secondary cold oxygen flow used to simulate the Earth’s atmosphere  composition. A converging-diverging nozzle is employed to expand the  hot mixture to a nominal supersonic Mach number equal to 3. The  general set-up used is illustrated in Fig. 1c,d. A detailed description of  the facility can be found elsewhere [22]. Some optical accesses allowed  a real time temperature monitoring. The specimen was mounted in front  of the nozzle exit (10 mm the distance) using a dedicated thermally  protected supporting mechanism [6]. According to Fig. 1d, the surface  temperature of the specimen was continuously measured (±1% instrumental accuracy) by digital two-color pyrometer (Infratherm ISQ5 and  IGAR6, Impac Electronic GmbH, Germany) at an acquisition rate of  100 Hz. An infrared (IR) thermo-camera (TC, Pyroview 512 N, DIAS  Infrared GmbH, Germany) operates in parallel and collects the signal  coming off the specimen according to Fig. 1c. The ISQ5 pyrometer exploits two overlapping infrared wavelength bands at 0.7-1.15 μm and  0.97-1.15 μm to measure the actual temperature from 1273 K up to  3273 K. The  (second)  IGAR6 pyrometer operates  in  the bands  1.5-1.6 μm and 2.0-2.5 μm to return the surface specimen temperature  in the range 523(cid:0) 2773 K. The fields of view of the ISQ5 and IGAR pyrometers are represented in Fig. 1d. The IR-TC covers a temperature  range from 873 to 3273 K and operates in the spectral range from 0.8 to  1.1 μm: its field of view of is highlighted in Fig. 1c. The procedure  employed to set the correct value of the spectral emittance ελ, on which  the IR-TC measurement is dependent, is reported elsewhere [19]. High  definition (HD) videos were on-line recorded by means of a video  camera (Flea3 1.3 M P Color USB3 Vision, resolution of 1328 × 1048  pixels, 25 frame/s grabbing rate).   2.3. Aero-heating testing   The four buttons were exposed to the supersonic plasma flow at  stepwise increasing specific enthalpy steps (see Table 1): the same test  sequence was executed on the four different buttons. The test sequence  starts setting a minimum current of 250 A which then is gradually  increased up to 600 A. The actual specific total enthalpy H0 was evaluated through a thermal balance [22]. The surface temperature of the  specimen was real-time monitored using the ISQ5 and IGAR6 pyrometers in combination with the IR-TC Pyroview 512 N, according to Fig. 1.  At the end of the most stressful enthalpy step, the torch power was  decreased gradually until facility shutdown, resulting in a stepwise  cooling of the buttons.  Upon cooling, an analytical balance (±0.1 mg of instrumental accuracy) and a micrometric caliper (± 1 μm) were used to determine,  respectively, mass change and size variation of the button after testing.  Prior to any sectioning, the as-exposed front surface of each button was  analyzed by XRD and FESEM-EDS.  To proceed for sectioning, each button was fully embedded in resin.   CorrosionScience178(2021)1090672\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 1. Lateral wall (a), front face (b) and dimensions of the flat button; arrangement of pyrometer ISQ5 or IGAR6 (c), and infrared (IR) thermo-camera (TC) with  video recorder FLEA (d): the colored portions of the button surface match diagnostics and their field of view.   Table 1  Current (I) and power (P) of the torch, specific total enthalpy (H0), cold wall  convective heat flux measured by a slug calorimeter (qCW) and stagnation point  oxygen partial pressure (pox).    Step   1  2  3  4  5  6  7  8   I, A   250  300  350  400  450  500  550  600   P, kW   12  15  18  21.5  25  29  33  36.5   H0, MJ/kg  6  8  10  12  14  16  18  20   qCW, MW/m2  0.95  1.1  1.5  2.0  2.5  3.0  3.6  4.2   pox, Pa  1240  1280  1340  1400  1460  1520  1580  1640    Then, a round slice, about 1 mm thickness and 17 mm diameter, was cut out from the top surface using a diamond blade 0.5 mm thick. This  procedure “refreshed” the button eventually ready for further aerothermal testing. The round slices were then cross-cut along a diameter  using the same diamond blade. The cross-sections so obtained were re embedded in resin, polished according to the method described in Section 2.1, and then analyzed by FESEM-EDS: Au coating of a few nm was  sputtered to prevent excessive electrostatic charging. Also elemental  mapping of selected areas of the oxide scale was done by FESEM-EDS.  Equilibrium compositions and reaction enthalpies of some reactions  were computed using the commercial package HSC Chemistry v.6.12  [23].   3. Results   The hot-pressing conditions were sufficient to bring the four ZrB2 - x  SiC - 5 Y2O3 compositions (x = 5, 10, 15 and 18.5 vol. %) to nearly full  density. In Fig. 2, the representative microstructure of each material  imaged through SE-FESEM revealed regular ZrB2 grains, SiC particulates  dispersed rather well  inter-granularly and residual glassy pockets.  Analogous analyses revealed measurable residual porosity in minor  amounts which was estimated through image analysis of FESEM micrographs (polished areas): the measured bulk densities (ρB) and the  estimated relative density (rd) are reported in Table 2.  Though  the  presence  of  oxides  often  affects   adversely   the   densification of non-oxide materials, in the present case the (intentional)  addition of 5 vol.% Y2O3 in the starting formulations and 2173 K of hot  pressing temperature did not impede to get fully dense materials.  Rather, Y2O3 interacted during hot-pressing with ZrB2 and SiC via their  native surface oxides, formed a transient liquid phase that, upon cooling,  remained concentrated  in discrete pockets which exhibit a glassy  appearance. An example is presented in Fig. 3a. XRD analyses identified  only ZrB2 and SiC as final crystalline phases present. However, FESEM EDS analyses of the glassy pockets assessed their main nature of a Y  borate. The uptake of B, as well as of Si and Zr but in minor amount, in  the glassy pockets is a robust evidence that, during hot pressing, Y2O3  got in touch with SiO2 and ZrO2-B2O3 (which are known being the native  oxides covering, respectively, the SiC and ZrB2 powder particle surfaces)  and reacted with them.  As for the arc-jet tests are concerned, during the test sequence, the  four buttons were subjected to all of the eight enthalpy steps of Table 1.  Steps from 1 to 7 had a duration of 30 s, while the number #8 lasted  120 s. After the most stressful step, the arc power was gradually  decreased until facility shutdown. The surface temperature vs time  (thermal) histories, measured by ISQ5 pyrometer according to the  configuration sketched in Fig. 1c, are plotted in Fig. 4. The visual  appearance of the four samples after test is presented in Fig. 5. All the  specimens gained a net mass, and increased the initial dimensions (see  Table 2).  Some interesting preliminary deductions can be argued. First, it is  clear that, whereas in the earliest segments of the power sequence the  surface temperatures do not vary appreciably, they tend to differ for  increasing enthalpy steps, until the maximum radiative equilibrium  temperatures, which appear be higher for decreasing SiC content, are  reached. Worth of mention is the behavior of specimen ZSY-05, which  experienced an extremely rapid temperature rise (recorded by ISQ5  pyrometer and IR-TC) when the surface temperature was approaching  2150 K, during step #7 (H0 = 18 MJ/kg). After reaching an almost  steady state, the temperature suddenly increased by more than 400 K.  The steady state equilibrium temperature would have probably exceeded 2600 K, but the power was increased to the last step before the  equilibrium was reached. The final peak temperature of the front surface  was around 2750 K. Another  interesting  feature was  that, at  the  maximum applied enthalpy H0, the surface temperature oscillates, with   CorrosionScience178(2021)1090673\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 2. SE-FESEM micrographs from a polished section of the as-hot-pressed materials: ZrB2 and SiC are the main phases and have, respectively, a prevailing grey and  dark contrast.   Table 2  Bulk density (rB) and relative density (rd) of the as-hot-pressed materials, absolute and relative mass gain (Δm) and thickness increase (Δt) of the button after  test, and thickness of the oxide scale after test (Δs).     rB, g/  cm3   5.78   5.74   5.62   5.51   rd, %   99.4   99.6   99.8   99.1   Δm,  mg   43.6   34.9   20.7   24.5   Δm/m0,  %   0.51   0.41   0.25   0.32   Δt,  μm   250   85   70   105   Δt/t0,  %   5.0   1.7   1.4   2.1   Δs,  μm   305   250   215   150    ZSY 05  ZSY 10  ZSY 15  ZSY 18   more pronounced amplitude variations for larger SiC contents in the  starting specimen.  The post-test visual inspection of the buttons assessed they have a  different appearance (Fig. 5). In particular, more SiC in the as-sintered  material led to get smoother front surfaces. In addition, a loss of   homogeneity is apparent in the ZSY-05 button, just that affected by the  steep temperature rise (see Figs. 4 and 5). For sake of consistency such a  phenomenon will be termed spontaneous temperature jump because it  occurs rapidly and at constant H0, and hereafter labeled TJ.  In Fig. 6, the overall cross-sections of the exposed specimens are  presented. It is apparent that the ZSY-05 button generated in-situ an  external oxide scale whose partial compactness (left half in Fig. 6a) was  severely altered (right half in Fig. 6a). The TJ, i.e. a temperature overshoot recorded by ISQ5 pyrometer and IR-TC, is a direct manifestation,  not the trigger, that a substantial change in the former thermal status  took place and a new one was equilibrated. More in specific, the ZSY-05  button presents irregular macro-granules which protrude for 50 up to  100 μm and, very likely, an oxide scale/un-oxidized bulk detachment  (right end in Fig. 6a). Some macro-granular structures on the front  surface of the ZSY-05 button are discernible also in Fig. 5. Another  important feature is the virtual absence of residual glass just on the  outermost front surface which directly faced the dissociated hot plasma.  New elements of understanding of the effects on the microstructure  altered by the exposure to the plasma were also provided by the XRD  analyses. In fact, Y became available during arc-jet test. The most   Fig. 3. Example by SE-FESEM of glassy pockets (1) taken from the as-hot-pressed ZSY-05 (a): also ZrB2 (2) and SiC (3) are numbered; b) FESEM-EDS spectrum (6 keV  of energy beam) from a glassy pocket: C signal is attributable to prolonged rastering during acquisition.   CorrosionScience178(2021)1090674\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             obvious effect was the stabilization of the zirconia phase composing the  external oxide scale, which prevented the transition to monoclinic  symmetry during cooling. XRD analyses on the exposed surfaces, besides  Y-stabilized zirconia with a tetragonal/cubic symmetry, identified also  Y2Si2O7. An example of the XRD analyses is presented in Fig. 7: in  particular, in Fig. 7b the contemporary presence of cubic and tetragonal  Y-stabilized zirconia is emphasized. FESEM-EDS analyses confirmed the  existence of the yttrium di-silicate upon the as-exposed oxidized front  surface, and of Y in the zirconia grains constituting the oxide scale.  Fig. 8 shows some thermographic frames taken during step #7 and  #8 of the test on sample ZSY-05, from the instant in which the TJ took  place until the end of the step #8. The frames are identified by numbers  1-8 and the corresponding instants are indicated on the diagram presented on the right of Fig. 8. This shows again the temperature history of  the sample as measured by pyrometer ISQ5, compared with two measurements made by the IR-TC, setting a directional spectral emissivity  equal to 0.75. The first, labeled as “center”, is the maximum temperature  detected on a round area in the central part of the sample front surface,  matching as close as possible the measurement spot of the ISQ5 pyrometer; the second, labeled as “max”, is the maximum temperature   Fig. 4. Surface temperature (T) versus time (t) using the ISQ5 pyrometer  during the test sequence of increasing H0 steps.   Fig. 5. Final post-test appearance of the four buttons.    Fig. 6. Polished cross-sections, about 17 mm wide, by BSE-FESEM of the tested buttons: upper and lower margins of the oxide scale (whose thickness is reported in  Table 2) are plotted.   Fig. 7. XRD pattern (background and Ka2 stripped away) of the exposed ZSY-15 button: 15-65 deg 2-theta range (a) and 33-37 2-theta range (b). Besides the Y stabilized tetragonal/cubic zirconia (YSZT/C), also a crystalline monoclinic yttrium di-silicate Y2Si2O7 was indexed. A clear example of the de-convoluted T/C  components of YSZT/C is shown in b).   CorrosionScience178(2021)1090675\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 8. Left: thermographic frames of button ZSY-05 during the most stressful enthalpy steps. Right: temperature history measured by ISQ5 pyrometer and IR-TC  (ελ = 0.75). For IR-TC two curves are shown: the maximum temperature on the center of the front surface (Center), in an area corresponding to the measurement  spot of the ISQ5 pyrometer; and the maximum temperature on the whole front surface (Max). The time instants corresponding to the TC frames are identified by  numbers 1-8 on the temperature history diagram.   detected on the entire front surface. Several observations can be made.  First, before the onset of the TJ (frame 1 of Fig. 8), the surface temperature was rather uniform, then a localized temperature increase  occurred on the front surface, spreading to the whole surface (but not to  the rear part of the material) when the new equilibrium temperature was  reached, at the beginning of the harshest enthalpy step (frame 4). During  the first half of the step #8 the three curves almost overlap, then for the  remaining time ranges (frames 6-8), the only “max” curve started  deviating and tended to reach values up to 3000 K. Such a temperature  drift was linked to the appearance of the protrusions on the top part of  the surface, with localized ultra-high-temperature spots.  Fig. 9 shows a sequence of IR-TC frames taken during step #8 of test  on the ZSY-15 button, in a time range characterized by oscillations in the  temperature history. It is possible to notice, on the bottom of the lateral  wall, an unsteady evolution of the irradiated power, resulting in a  change of the surface color. An analogous behavior was observed also  for the remaining buttons. A similar phenomenon was observed by  Monteverde et al. [19], who called it “waves of radiance” and correlated  it to the transport of a liquid glassy phase from the front surface along  the lateral walls of the specimen by the shear stresses induced by the  supersonic flow.  Fig. 10 shows the (directional) spectral emissivity (ελ) of the four  specimens versus surface temperature, obtained by matching the measurement of the ISQ5 pyrometer and IR thermo-camera, in the near  infrared (NIR) wavelength band. For all samples, the ελ values tend to  increase at the beginning of the test, reach a maximum that ranges between 0.8 and 0.9, at temperatures between 1600 and 1700 K, then they  begin to decrease converging to higher values for the compositions  richer in SiC. After TJ, ελ of the ZSY-05 button (the only that exceeded  2400 K) increases again up to 0.75 at the maximum temperature of  2750 K. One could argue that the temperature and emissivity measurements after TJ are less reliable due to the presence of protrusions,  leading to inhomogeneity in temperature distribution on the sample   Fig. 10. Directional spectral emissivity (ελ) versus maximum surface temperature (T).   front surface. However, a spectral emissivity trend comparable with that  of the ZSY-05 button has already been observed in a previous work [21]  regarding test in similar conditions on fiber-reinforced UHTCs based on  ZrB2-SiC-Y2O3, also experiencing a TJ. In that case, no significant protrusions were observed on the samples front surface, which appeared  homogeneous and compact. This is a first reason to infer that also the  present temperature and emissivity measurements may not be affected  by the specific presence of large protrusions. Moreover, looking at Fig. 8,  it can be seen that in the first half of step #8, when the final equilibrium  temperature of 2750 K had already been reached, the front surface  temperature was uniform and the measurements of pyrometer and  thermo-camera almost perfectly overlapped. This suggests that the set  value of ελ = 0.75 was already representative of the actual (spectral)  emissivity of the surface even before the catastrophic bursting of large  protrusions and defects in the second half of the step (frames 6-8).  Fig. 11 shows the surface temperature profiles along the axis of the   Fig. 9.  IR-TC frames grabbed during step #8 showing waves of radiance along the lateral wall of ZSY-15 button.    CorrosionScience178(2021)1090676\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             button, measured at the same time instant by the IR-TC during step #8.  As already highlighted above, the substantial temperature rise related to  the TJ appears to be confined on the front surface of specimen ZSY-05,  with a steep temperature drop alongside the axial coordinate of the  button. A similar, but less obvious trend, can be observed for the  remaining buttons and the magnitude of the temperature drop decreases  for increasing SiC content, suggesting a progressive increase in the  ability of the oxide scale to conduct heat to the bulk material. The formation of oxidized areas on  the side surfaces, which  is actually  confirmed also by SEM analyses (see Figs. 6-8), can of course have some  impact on the thermal conductivity-driven heat transfer to the rear part  of the samples. However, the temperature profiles reported in Fig. 11 are  the only directly available data to study the thermal axial distribution.  Regardless of the accuracy of those profiles for a precise reconstruction  of the heat transfer to the material, the variation, with SiC content, in  their slope in the front area provides a clear, although only semi quantitative, insight on the reduction of the thermal conductivity in  that region of the sample, supporting the discussion that will be presented in next section.  The resulting thickness of the oxide scale (see Fig. 6) is visibly  affected by the starting SiC content. Larger SiC amounts favored the  creation of thinner and more compact oxide scale: if the oxide scale of  the ZSY-05 button is thicker (i.e. about 300 μm) and partially detached,  in the ZSY-18 button it remained much more compact and thinner (i.e.  about 150 μm), and adherent to the un-oxidized substrate.  The FESEM-EDS analyses of the in-situ generated external oxide scale  provided an additional insight into the internal zones affected by a severe oxidation attack. Two important shared features, often observed by  other authors, were the lack of a SiC-depleted region underlying the  oxide scale and the prevailing columnar shape of the zirconia grains  grown during exposure [24]. If the columnar shape of the zirconia grains  is apparent along the lateral walls of the exposed button, the elongated  shape (of the zirconia grains) turns into a more regular one especially  close to the outermost part of the oxide layer where temperature reached  the hottest values. The change in shape and size of the zirconia grains,  which in the early stages of the test grew acquiring a columnar direction,  is a consequence of their sintering. The directional growth of zirconia  grains was fed by the incoming oxygen which diffuse inward through  paths created by the capillarity-driven rise of the fluid glassy oxide. The  other gaseous oxidation by-products originated at the oxide/inner bulk  interface such as SiO(g), CO(g) and volatile boron oxides counter-diffuse  due to concentration gradient exploiting the cited channels to reach the  exterior.  Such a fluid glassy oxide finds origin from the oxidation products of  the passively oxidized ZrB2 and SiC, according to the well-known   Fig. 11. Temperature axial profiles of the four buttons at the same instant  (t = 300 s) during step #8, measured by the IR-TC (the correct value of ελ was  set according to the diagram of Fig. 10).   exothermic reactions    ZrB2(s)+ 5/2 O2(g) = ZrO2(s) + B2O3(l,g)                                            SiC(s) + 3/2 O2(g)= SiO2(s,l) + CO(g)                                                (1)    (2)   Fed also by the residual Y borate glassy pockets, an only oxidation  product tends to form, up-taking Si, Zr, B and Y in only one oxide matrix:  such an oxide matrix is generally defined borosilicate glass (BSG). The  BSG, depending on the pressure and temperature it actually encounters  during the exposure, can vary its own composition substantially. To  mention the most evident change, it begins losing the boron oxide  component when the massive volatilization of the silica component has  not started yet. The channels grown inside the oxide scale acted as  getaway for the outflowing glass and other gaseous species and as an  access gate for the incoming oxygen. An example of the channels as well  as localized irregular ramifications is shown in Fig. 12. Part of the discussion will address this point.  The elemental mapping of Si as well as other key species such as Y, Zr  and O throughout the oxide scale is very important (Figs. 13 and 14)  because  it discloses how the BSG glass was transported outward,  providing or not protection to the specimen. The response to aero thermal heating was different. During the exposure of the ZSY-05 button, the entire oxide scale was virtually depleted of Si (Fig. 13). However, in the case of the ZSY-18 button, residual Si remained inside the  oxide scale localized in small glassy pools, whose presence and distribution are discernible in Fig. 14. The maps of Figs. 13 and 14 depict the  elemental distribution of the button poorest (Fig. 13) or richest in SiC  (Fig. 14). Analogous elemental analyses were carried out in the button  ZSY-10 and ZSY-15, and intermediate distributions (not shown) were  similarly found.  In the ZSY-05 button, a Si-free strip of several tenths of micron survived, sandwiched between the oxide/substrate interface and the zone  rich of residual glassy pools. Such an inner strip was interested by the  diffusion of Si transported via SiO(g). As for the lack of SiC-depleted  region underlying the bottom of the oxide scale, this can be explained  in terms of the oxygen partial pressure vs temperature at the oxide/bulk  interface which did not reach the threshold to switch into the active  oxidation domain. The reasoning over the SiO(g) outward diffusion and  its re-condensation will be reclaimed in the next section.   4. Discussion   With the support of the experimental determinations just reported in  the “Results” section, it was verified that the variation of ZrB2 to SiC  volume ratio induced measurable different effects on the aero-thermal  behavior in high-enthalpy supersonic airflows. One of the most striking features was the difference in the maximum equilibrium temperature when exposed to the same airflow conditions. According to Figs. 10  and 11, this is related to transient changes of the surface (spectral)  emissivity, at least in material ZSY-10, ZSY-15 and ZSY-18 where no  obvious TJ occurred, and very likely to a decay of the capability of the  oxide scale to conduct heat.  The observed correlation of SiC concentration vs the steady-state  temperature is in agreement with what found by Hu et al. [25], who  tested ZrB2-based specimens with a SiC amount ranging from 10 to  30 vol. % in a high-enthalpy subsonic wind tunnel. They also found a  reduction in the equilibrium temperature for increasing SiC content,  associated with an increasing amount of glass retained in the sample  after testing. In fact, it is largely accepted that more SiC promotes a  better physical barrier against oxidation thanks to the presence of a  borosilicate glass (BSG) [9,10,26]. Actually, the “durability” of the BSG,  i.e. its capacity to remain not only adherent to the external surface but to  survive also inside the oxide scale, is vital to continue having an effective  barrier to oxidation: in respect to the four compositions studied, this did  take place only in part. Due to the temperatures reached during the  present test campaign, on the front surface of the buttons, the BSG got   CorrosionScience178(2021)1090677\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 12. BSE-FESEM micrograph (polished section) from the front edge surface of ZSY-15 button (a): the directionality of the channels is evident; b) ramifications in  the box (a) are magnified.   Fig. 13. BSE-FESEM micrographs from front areas detached (a), or not (b) by the un-oxidized substrate of the ZSY-05 button: residual resin (*) is indicated.  Elemental maps of Si, O, Y and Zr derive from the selected boxes.   lost. In fact, according to the post-test analysis by FESEM-EDS, the  exposed surfaces showed sparse residual glassy pools, unable alone to  protect effectively from the oxidation attack. An as much precious part  of the BSG survived inside the oxide scale, in larger quantities for more  SiC available in the starting composition. If the outermost glassy layer  was basically lost during the most stressful stages of the test sequence,  the underlying oxide scale was able to hold more and more BSG as long  as the thermal insulation of the oxide scale worked in an effective  manner. In Figs. 13 and 14, the concentration profile of Si reflects that of  the residual BSG in button ZSY-05 and ZSY-18: larger quantities of BSG  survived in the oxide scale of the SiC richest composition. Nevertheless,  the lack of BSG noticed inside the Si-free strip lying above the oxide/bulk interface (see Fig. 14) was interpreted in this manner. SiO(g), and  not liquid BSG, forms at the oxidation front, diffuses through the open  channels outward, leaving voids and porosities behind. If such physical  discontinuities are continuously replenished by fresh BSG rising up from   the interior and re-condensing for the increased oxygen partial pressure,  the oxide scale is still able to afford protection against oxidation and  draw heat away from the stagnation points. The concentration profile of  the residual Si in the ZSY-18 specimen tells us that, for the duration of  the test, the replenishment seamlessly worked so that the feeding of  protective BSG never stopped. This is more evident along the lateral  walls of the button which were affected by less harsh conditions of heat  flux and temperature compared to the front one (see Fig. 11). In fact, the  FESEM-EDS mapping revealed the presence of residual BSG close to the  deepest oxidation front, with increasing residual amount of the same  moving toward the exterior.  Such a feeding mechanism of BSG was not effective in specimen ZSY 05 due to a native lack of Si, drastically shrinking the “incubation” time  necessary to deprive the oxide scale of the residual BSG. The concept of  an incubation time preceding a TJ was already mentioned by other  authors which correlated  this phenomenon  to  the  rapid  loss of   CorrosionScience178(2021)1090678\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 14. BSE-FESEM micrographs from the front surface (a) or the lateral wall (b) of the ZSY-18 button: elemental maps of Si, O, Y and Zr derive from selected boxes.    protection no more afforded by external protective glasses [12,20,21].  The depletion of BSG inside and outside the zirconia oxide scale was  pre-condition to push the ZSY-05 button to interact differently with the  flow-field, disclosing in a measurable manner a new energy balance  through a rapid spontaneous rise of the front surface temperature, e.g  TJ.  It should be reminded  that reactions  (1) and  (2) never stop  happening until oxygen diffuses inward and reaches the un-oxidized  bulk: they are exothermic and in this sense they continue to release  extra heat to the system. On the contrary, the loss of the BSG implies  endothermic reactions. As for the silica component of the BSG is concerned, some reactions like    SiC (s) + 2 SiO2(l) = 3 SiO(g) + CO(g)ΔHR (2200 K) =685.3 kJ/mol     (3)    SiO2(l) + CO(g) = SiO(g) + CO2(g)ΔHR (2200 K) =489.8 kJ/mol            SiO2(l) = SiO2(g) ΔHR (2200 K) = 568.9 kJ/mol                                     (4)    (5)   can simultaneously occur, leading the specimen to lose mass. The  equilibria calculations of ideal phases, although the volatilization of the  BSG is rather complex, anyhow remain a useful tool to draw trends,  some of them are plotted in Figs. 15 and 16. The calculated equilibria  point out how temperature, pressure and carbon (carried by SiC or CO)  affect the volatilization rate of BSG. The initial composition of one mole   BSG was approximated by adding 0.2 mol B2O3 to 0.8 mol SiO2, in one  mole of air (0.8 N2 + 0.2 O2). The activities of phases were all set equal  to 1, while the equilibrium pressure was let vary. In this regard, the  outputs in Fig. 15 do not take into the effect of the continuous removal of  gaseous products by means of the vacuum pumping system which shifts  the equilibrium constant toward lower temperatures: in this sense the  ranges of temperature and pressure along which SiO2 and B2O3, the two  main constituents of the BSG, begin to volatilize are, respectively, over and under-estimated. In any case, a reaction enthalpy ΔHR in the range  of 600(cid:0) 700 kJ/mol was estimated for the volatilization of BSG. The  presence of C (carried by SiC or CO) shifts the equilibrium amounts of  the most stable phases composing the BSG towards lower temperatures.  It was paid attention to the fact that, when the surface gets depleted  of BSG, a quote of the incoming heat flux, till now trapped to feed the  BSG volatilization, returns available to provide an extra heating. Actually, the overall amount of the consumed SiC, i.e. that formerly present  in the volume presently occupied by the oxide scale, is about 3 × 10(cid:0) 4  and 5 × 10(cid:0) 4 mole, respectively, in ZSY-05 and ZSY-18 specimen, but so  limited that any extra heat eventually released to the system (only a few  W) has to be considered of marginal entity.  The different behavior of various compositions containing different  initial contents of SiC, and ultimately the TJ, can be likely explained  according to the above described oxidation process, and taking into   Fig. 15. Equilibrium amounts (Q) versus temperature (T) in 1 mol air at overall pressure of 0.01 bar, 0.8 SiO2 + 0.2 B2O3 with/without 0.5 C starting composition: a)  most stable solid/liquid (s,l) phases; b) gases (g). Transition from solid to liquid SiO2 occurred at 2000 K.   CorrosionScience178(2021)1090679\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 16. Equilibrium amounts (Q) versus pressure (P) in 1 mol air at 2200 K of equilibrium temperature, 0.8 SiO2 + 0.2 B2O3 with 0.5 C starting composition: a) main  stable liquid (l) and gaseous (g) phases; b) minority gases. The typical range of pressure at the stagnation points varies from 5 to 9 kPa.   account a combination of factors. The TJ onset appeared taking place in  a localized spot which grows in size versus time and tends to affect wider  portions of the only front surface exposed to the plasma torch. The  experimental determinations completed using the IGAR6 pyrometer and  IR-TC (see Fig. 8) excluded the occurrence of anomalous temperature  rises along the lateral walls of the button.  The hypothesis here proposed is that the TJ occurred in the specimen  which first lost the protection of the BSG not only lying on the external  surface (irreparably wiped away by shear forces) but mainly present  inside the oxide scale. It is essential restate the equally important  twofold role of the BSG as “filler” to keep glued the zirconia grains of the  oxide scale and barrier against oxidation. If the protective capability is  obvious, the capacity of the BSG to retard the onset of defects such as  voids or even detachment is as essential as elusive to be assessed. As far  as the oxide scale is concerned, the localized growth of irregular ramifications (see Fig. 12) which can lead to the formation of the macro granular protrusions, and the deteriorated compactness due to the creation of defects are a couple of  factors which may enhance  the  convective component of the heat flux accompanied by a reduced capacity to draw heat away. Such a scenario was thought being the trigger  at the basis of the spontaneous temperature jump. Actually, aero-heating  of ZSY-05 specimen was not stopped but proceeded through the most  stressful last step. It follows that many microstructure features, more  pertinent to the scenario above described, may have been irreparably  changed. In fact, after the TJ, the IR-TC outputs confirmed that some  front areas of the ZSY-05 button reached temperatures so high to induce  a partial melting of the zirconia oxide scale directly exposed to the  harshest heat flux condition (see Fig. 8). However, looking at the lateral  walls of the same button which experienced less severe heating conditions and lower surface temperature, the formation of irregular ramifications and voids is apparent. Also, a detachment of the oxide scale is  clear, though it cannot be excluded that it occurred, fully or partly,  during cooling down: the IR frames of Fig. 8 suggest that the area which  underwent the oxide detachment corresponds to the spot of the further  temperature increase measured by the IR-TC in the second half of step  #8. Moreover, even the ZSY-15 button (richer in SiC compared to ZSY 05) presents a longitudinal subtle opening of the oxide scale (see  Fig. 6c), although no apparent TJs have been measured and no catastrophic damage to the oxide layer was observed.  It was paid great attention to an oxide scale/substrate detachment; as  a very critical event it may occur when specific mechanisms proceed  forward acting all together uncontrolled during the exposure at very  high temperature: the disruptive escape of gases out to the exterior,  the “embrittlement” of the oxide scale depleted of BSG; the tendency of  the oxide scale to dilate more than the underlying un-oxidized substrate.  Depending on  the extent of  the “oxide scale embrittlement”,  the  macroscopic effect can vary from a partial exfoliation of the oxide scale  (observed in specimen ZSY-15) to its detachment from the un-oxidized   substrate (observed in specimen ZSY-05). The oxide scale of the ZSY 18 button, based on the resulting features imaged via FESEM, did not  suffer appreciable damage because it retained enough BSG during  exposure: this favorable combination allowed the ZSY-18 material to  endure very well the hostile aero-heating conditions. In general, the  temperature axial profiles along the four samples, depicted in Fig. 11,  make possible to infer that the ability of the materials outer (oxide)  layers to conduct heat during the most stressful test stage was reduced  accordingly with the capability of retaining BSG, thanks to the presence  of higher SiC concentration in the initial composition. This interpretation finds consensus with the relevant literature investigating the ultra high-temperature heat transfer of HfB2and ZrB2-based UHTCs [21,27,  28], where porous zirconia layers were even proposed as effective  insulating thermal barrier coatings for high-temperature aerospace  components [29,30].  Indeed, the effective thermal conductivity of the oxide scale can  substantially affect the distribution of the incoming convective heat flux  among the conductive and radiative contributions, resulting in highly  different radiative equilibrium temperatures on the exposed surface.  Referring to the conditions measured on sample ZSY-05 during step #7,  a front surface temperature of 2200 K before TJ and of 2600 K after TJ  were here assumed. By numerical models described elsewhere [3,20],  the distribution of the oncoming convective hot wall heat flux among its  radiative and conductive components has been estimated. Fig. 17 shows  the effect of a 300-μm thick oxide scale thermal conductivity (in the  range 0.1-2.5 W/m‧K) on the conductive heat flux, together with horizontal lines corresponding to the values of radiative heat flux before and  after TJ (for a fixed surface emissivity). It can be seen that, to ensure   Fig. 17. Variation of the conductive heat flux (q˙cond) versus the thermal conductivity of the oxide scale (kox). Horizontal lines corresponding to the values of  the radiative heat flux before and after TJ (q˙rad,pre and q˙rad,post respectively) are  also shown.   CorrosionScience178(2021)10906710\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             global equilibrium, the substantially increased surface radiation due to  the higher surface temperature can be balanced by a reduced capability  to dissipate heat by conduction. Specifically, it has been estimated that,  in the conditions before TJ, assuming the presence of an oxide layer with  a thermal conductivity of 2.5 W/m‧K (reasonable for a zirconia scale  whose internal voids and porosities are filled by liquid silica-based glass  [31,32]), the total hot-wall convective heat flux consists of 35 % radiative dissipation towards the environment and 65 % conduction to the  inner layers of the material. Considering, for the conditions after TJ, a  drastically  reduced  effective  thermal  conductivity of 0.1 W/m‧K  (resulting from the interruption of the BSG replenishment and the corresponding creation of voids, porosities [33] or even a detachment), the  total heat flux redistributes to 70 % radiation and 30 % conduction.  To recap, the new thermal status just after a TJ reflects the energy re balancing of the system exposed to the same freestream energy input but  affected by a modified capability to drain away heat in excess from the  stagnation points and thus to re-radiate it: the system “spontaneously”  equilibrates a new status where the radiative temperature is globally  higher.   4.1. Role of Y and its effects on the aero-heating resistance capabilities   The expected impact of Y as stabilizer of the zirconia grains was  fulfilled. In this respect, according to the Y concentration profile of ZSY 05 specimen in Fig. 13, the quantity of Y inside the zirconia grains increases outward and reaches its peak of about 12 at% (i.e. 20 % mole) in  those grains just facing the external atmosphere. Based on the XRD  analyses which indexed a part of the stabilized zirconia with a cubic  symmetry, it is plausible infer that cubic Y-stabilized zirconia resides the  outermost regions of the oxide scale because it holds more Y and is more  prone be formed at very high temperature. However, in the ZSY-18  button,  thanks  to reduced surface  temperature,  the concentration  gradient of Y resulted less pronounced.  Independently of the starting composition, part of the available Y  reacted and formed Y2Si2O7 crystals lying on the external front surface  (see Fig. 7). Another Y-linked feature was found inside the oxide scale: a  Y-based oxide. Very  likely, because of miscibility  limitations,  it   separated from the BSG and crystallized assuming a brain-like layered  structure (see Fig. 18).  The elemental mapping of Fig. 19 clearly exhibits  the  interpenetrated arrangement of the Y-based oxide separated by the residual  BSG. Interestingly, the Zr map showed a not negligible intensity of  counts inside the brain-like feature. Actually, the intimate interpenetration of such a Y-based oxide with the residual glass puts the FESEM EDS technique very close to its limits and impeded to discriminate its  actual chemical composition. Semi-quantitative analyses (see Fig. 18d)  make very unlikely we are still in presence of Y2Si2O7, which instead was  clearly identified by XRD and FESEM-EDS on the exposed surface. More  likely, an (Y,Zr)2O3 separated from the BSG oversaturated of Y.   5. Conclusions   A focused testing campaign explored the effects of varying the SiC  concentration on the aero-thermal resistance capabilities of ZrB2-based  UHTCs. Four ZrB2 - x SiC - 5 Y2O3 compositions, x = 5, 10, 15 and  18.5 vol. %, were fully densified by hot pressing. Lab-scale flat buttons  (17 mm diameter) were fabricated and then tested in an arc-jet supersonic wind tunnel environment increasing the specific total enthalpy up  to 20 MJ/kg. It was found that, applying the same freestream conditions,  as SiC content  increases,  the maximum  steady-state  temperature  reached on  the button surface decreases. Microstructure analyses  assessed the formation of Y-stabilized tetragonal/cubic zirconia-based  oxide  scales which  (excluded  the composition x = 5 vol. %) also  contain residual contents of a borosilicate glass. The ZrB2-18.5SiC 5Y2O3 composition endured rather well the aero-heating conditions up  to 20 MJ/kg specific total enthalpy: the overall oxide scale, about 150  μm thick, resulted compact and smooth. The ZrB2-5SiC-5Y2O3 behaved  differently because it was the only affected by a sudden rise in temperature in the order of 400(cid:0) 500 K under freestream conditions at  18 MJ/kg of specific total enthalpy. The temperature jump affected only  the front surface directly facing the hot plasma while the lateral walls of  the button remained cooler. The origin of such phenomenon was related  to the complete loss of protective glass, and was then explained in terms  of a substantial deteriorated capacity of the oxide scale to drag heat   Fig. 18. BSE-FESEM micrographs (a,b) from ZSY-18 button: residual silica glass (1), Y-stabilized zirconia (2), and the separated Y-based oxide are numbered;  corresponding FESEM-EDS spectra (c,d) are also shown (6 keV of energy beam).   CorrosionScience178(2021)10906711\\x0c', 'S. Mungiguerra et al.                                                                                                                                                                                                                             Fig. 19. BSE-FESEM micrograph (a) from a selected area of ZSY-18 button showing the main phases composing the oxide scale: Y-stabilized zirconia (1), BSG (2) and  the brain-like silica/Y disilicate (3); elemental maps of Si, O, Y and Zr were generated from the box (15 keV energy beam).   delivered to the surface. The oxide scale was moreover supposed to  suffer a partial detachment from the un-oxidized substrate, creating a  new physical discontinuity. The new thermal status just after the jump  represented the response of the material to re-radiate the heat in excess,  no more effectively conducted away from the stagnation points. Additional tests which address longer durations at specific total enthalpies  below 18 MJ/kg might better disclose if prolonged incubation times and  heat  load  thresholds are strongly correlated  to provoke TJ. Also  computational fluid dynamic simulations will address the interpretation  of the effects of the flow conditions (including shear stresses) and material properties (surface catalycity, emissivity, thermal conductivity) on  the TJ onset.   Data availability   The raw/processed data required to reproduce these findings cannot  be shared at this time due to technical or time limitations.   CRediT authorship contribution statement   Stefano Mungiguerra: Conceptualization, Methodology, Validation, Formal analysis, Investigation, Data curation, Writing original  draft, Visualization. Anselmo Cecere: Investigation, Writing review &  editing. Raffaele Savino: Conceptualization, Methodology, Supervision, Funding acquisition. Federico Saraga: Methodology, Investigation.  Fr´ed´eric  Monteverde:  Conceptualization,  Methodology,  Investigation, Formal analysis, Resources, Visualization, Writing   original draft. Diletta Sciti: Supervision, Funding acquisition, Project  administration.   Declaration of Competing Interest   The authors report no declarations of interest.   Acknowledgements   The C3HARME research project has received funding by the European Union’s Horizon2020 research and innovation programme under  the Grant Agreement n  685594.   References   [1] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for  extreme environments, Scr. Mater. 129 (2017) 94-99, https://doi.org/10.1016/j.  scriptamat.2016.10.018.  [2] R. Savino, S. Mungiguerra, G.D. Di Martino, Testing ultra-high-temperature  ceramics for thermal protection and rocket applications, Adv. Appl. Ceram. 117  (2018) s9-s18, https://doi.org/10.1080/17436753.2018.1509175.  L. Silvestroni, S. Mungiguerra, D. Sciti, G.D. Di Martino, R. 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Therm. Spray Technol. 23 (2014) 849-859, https://doi.org/10.1007/  s11666-014-0082-5.  [30] G. Ouyang, M.F. Besser, M.J. Kramer, M. Akinc, P.K. Ray, Designing oxidation  resistant ultra-high temperature ceramics through the development of an adherent  native thermal barrier, J. Alloys Compd. 790 (2019) 1119-1126, https://doi.org/  10.1016/j.jallcom.2019.03.250.  S. Elhadj, S.T. Yang, M.J. Matthews, D.J. Cooke, J.D. Bude, M. Johnson, M. Feit,  V. Draggoo, S.E. Bisson, in: G.J. Exarhos, V.E. Gruzdev, D. Ristau, M.J. Soileau, C.  J. Stolz (Eds.), High Temperature Thermographic Measurements of Laser Heated  Silica, 2009, p. 750419, https://doi.org/10.1117/12.836985.  S. Kasap, J. M´alek, R. Svoboda, Thermal properties and thermal analysis:  fundamentals, experimental techniques and applications. Springer Handb.  Electron. Photonic Mater., Springer International Publishing, Cham, 2017, p. 1,  https://doi.org/10.1007/978-3-319-48933-9_19.  B. Nait-Ali, K. Haberko, H. Vesteghem, J. 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},{
  "_id": 97,
  "PDF": "Improved Oxidation Resistance of Zirconium Carbide at 1500°C by Lanthanum Hexaboride Additions.pdf",
  "Text": "['Improved Oxidation Resistance of Zirconium Carbide at 1500°C by  Lanthanum Hexaboride Additions  Liyou Zhao, Dechang Jia*,†  Xiaoming Duan, Zhihua Yang, and Yu Zhou  Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin, 150001, China  Addition of LaB6 is adopted to improve the oxidation resistance of ZrC at 1500°C. Mixed powder of ZrC-25 vol% LaB6 is reactively hot pressed at 1900°C for 30 min under vacuum  with  an  applied  pressure  of  25 MPa. The LaB6 ZrC to form ZrB2 and a layered La-containing phase. ZrB2 improves the oxidation resistance of ZrC in static air. The  reacts with  La-containing  phase  is  beneficial  to  increasing  the  relative  density of oxide  scale during oxidation and in enhancing the  oxide scale stability during exposure to thermal cycles.  I.  Introduction  Z IRCONIUM carbide properties such as high melting temperature, high thermal and electrical conductivity, and high fracture strength.1-3 It is a potential ultrahigh-temperature ceramic (UHTC) for aerospace applications.4,5  (ZrC)  shows  a  number  of  excellent  Oxidation resistance  is a major  issue  in the development  of UHTCs. The ZrC has a poor high-temperature  chemical  stability  in  air, which yields a porous oxide scale during oxidation.6,7 Until now, the reports on enhancing oxidation resistance of ZrC are scarce. He et al. 8,9  and  non-protective  prepared ternary  and quaternary  carbides by  incorporation  of Al or Al with Si  into ZrC. The obtained compounds show  better oxidation resistance than ZrC due to the formation of  some protective products, but  the oxide  scale  is  still porous  because of the inability to sinter and the formation of high vapor pressure gaseous products. Recently, Zhang et al.10,11  introduced 4 mol% WC into ZrB2 matrix, which promotes formation of a liquid phase during oxidation and conse quently results  in liquid-phase  sintering of ZrO2 of oxide scale and  scale. This  increases  the  relative  density  therefore  decreases  the  rate  of  oxygen  transport. ZrB2-4 mol% WC performs much better oxidation resistance than monolithic 1500°C and 1600°C.  ZrB2 strategy  at  It  is  a  novel  and  promising  to improve  the oxidation resistance of UHTCs by  forming  liquid during oxidation and in turn enhancing  the  protective effect of oxide scale.  In a previous work, we improved the fracture toughness of in situ  ZrC by addition of LaB6, which is attributed to the formation of a layered La-containing compound.12  In this  article, we report  the enhancement of oxidation resistance of  ZrC by LaB6 addition. It provides a new method for enhancing oxidation resistance of ZrC.  II.  Experimental Procedure  size 2.1 lm, The starting powders were ZrC (mean particle >98% purity, Changsha Wing High High-Tech New Materials Co., Ltd., Changsha, China), ZrB2 (D50 = 2.5 lm, 99% purity, Northwest Institute for Non-ferrous Metal Research, (\\x00300 mesh, >99% purity, NorthXi’an, China), and LaB6 west Institute for Non-ferrous Metal Research). Mixed pow der of ZrC-25 vol% LaB6 was reactively hot pressed in boron nitride-coated graphite die at 1900°C for 30 min under vacuum  with an applied pressure of 25 MPa. The obtained sample is  referred as CBL in this study. Monolithic ZrC and ZrC-40 vol  % ZrB2 composite were also prepared. The sintering temperature was 2000°C, holding time 60 min, and applied pressure 25 MPa. Cylindrical coupons with dimension of Φ12.7 9 4 mm  were cut from the hot-pressed plates. Static oxidation test was  carried out by exposing coupons to stagnant air in a box furnace at 1500°C for 2 h or 4 h. The heating rate was 15°C/min. Cyclic oxidation test was also conducted. The coupons were exposed to stagnant air at 1500°C for 15 min and then removed quickly  from the furnace, and left  for cooling in air  for 10 min. The  total time in furnace was 2 h. After oxidation, the oxide scale  was ground off and the remained thickness of coupons was mea sured to evaluate the oxidation degree of samples. The phase  composition of oxide scale was identified using X-ray diffrac(XRD, CuKa radiation, D/max-2200VPC, Rigaku,  tometer  Tokyo, Japan). The surface and cross-section of oxidized sam ples were observed by scanning electron microscope  (SEM,  FEI  Quanta  200F,  Eindhoven,  the Netherlands),  which  is  equipped with  an  energy-dispersive  spectroscopy  (EDS,  EDAX, NJ) system.  III.  Results and Discussion  The obtained composite using ZrC and LaB6 as als is composed of three phases, namely, ZrC, ZrB2, and a [Fig. 1(a)]. Element La mainly distributes in  raw materi layered phase  the layered phase [Fig. 1(b)]. The layered phase also contains elements B and C.12 After static oxidation at 1500°C for 4 h, CBL forms a light yellow oxide scale, and no cracking of oxide scale or edge  cracking of  sample is observed [Fig. 2(a)]. Some bubbles are  detected on the  surface of  sample,  suggesting formation of  liquid and gaseous phases during oxidation. The  remained  thickness of coupon is 2.96 mm. ZrC oxidizes catastrophically in air at 1500°C. The oxide scale significantly spalls from the  substrate [Fig. 2(b)], and the remained thickness of coupon is  1.67 mm. ZrC-40 vol% ZrB2 and edge cracking after static oxidation at  sample  exhibits severe surface 1500°C for  4 h  [Fig. 2(c)], but  the remained thickness of coupon is the thick est, about 3.28 mm. After 2 h cyclic oxidation, the oxide scale  of CBL appears  dense  and  adherent  [Fig. 3(a)],  and  the  remained thickness of  coupon is  3.32 mm;  in contrast,  the  oxide  scale  of  ZrC-40 vol% ZrB2 [Fig. 3(b)], and the remained thickness of coupon is 3.16 mm.  is  damaged  seriously  B. Fahrenholtz—contributing editor  Manuscript No. 29625. Received April 23, 2011; approved August 05, 2011.  This work was  supported by the Program for Changjiang Scholars and Innovative  Research Team in University.  *Member, The American Ceramic Society.  †  Author  to whom correspondence should be addressed. e-mail: dechangjia@yahoo.  com.cn  3648  J. Am. Ceram. Soc., 94 [11] 3648-3650 (2011)  DOI: 10.1111/j.1551-2916.2011.04830.x  © 2011 The American Ceramic Society  Journal  \\x0c', 'Figure 4 shows the surface SEM micrograph of different samples after static oxidation at 1500°C for 2 h. The surface  of CBL appears  relatively dense. The dark phase  in back scattered electron (BSE)  image is ZrO2. Glassy phase is also observed, which mainly distributes at the boundaries of ZrO2 at the boundaries. EDS  grain clusters  and seals  the  cracks  analysis  (not  shown)  indicates  that  the glassy phase contains  elements La, O, and possibly B. The oxide  scale on ZrC is  cracked and porous while the scale on ZrC-40 vol% ZrB2 porous.  is  Introduced LaB6 reacts with ZrC to form ZrB2 and a layered La-containing phase. After static oxidation test, the  remained coupon of CBL is  thicker  than that of ZrC, but  thinner  than that of ZrC-40 vol% ZrB2. So, CBL has better static oxidation resistance than ZrC, but inferior static oxida tion resistance compared with ZrC-40 vol% ZrB2. Assuming that all B in the introduced LaB6 is converted to ZrB2, the of ZrC to ZrB2 in CBL is about 3:2. the enhancement of oxidation resistance of ZrC in static air  volume  ratio  So,  after LaB6 addition is attributed to the  formation of ZrB2.  However, we note that  the oxide scale on ZrC-40 vol% ZrB2 cracked macroscopically after cooling while the scale on  is  CBL is dense and adherent. Microstructure analysis  reveals  that CBL has  denser  oxidized  surface than ZrC-40 vol% 1500°C,  ZrB2. After coupon of CBL is  2 h  cyclic  oxidation  at  the  remained  thicker  than that of ZrC-40 vol% ZrB2, cyclic oxidation resistance of CBL. So,  indicating the better  the La-containing phase  in CBL is beneficial  to increasing  the  relative  density  of  oxide  scale  during  oxidation  and  enhancing the oxide scale stability during cyclic exposures.  Figure 5 shows the XRD patterns of oxide scale on CBL static oxidation at 1500°C for 2 h and 4 h. The oxide formed at 1500°C for 2 h is composed of monoclinic  after  scale  ZrO2 No.  (JCPDS No.  24-1165), orthorhombic LaBO3 hexagonal La2O3 (JCPDS No. Besides orthorhombic LaBO3, another crystal form of LaBO3 (JCPDS No. 13-571) is found in the pattern of the oxide 1500°C for  (JCPDS  12-762),  and  83-1345).  scale  formed  at  4 h. The  high  background  in  Fig. 5(b)  is  consistent with the presence of  glassy phase  in  the oxide scale. By far,  it  is difficult  for us to clarify the oxi dation mechanism of La2O3-B2O3 important role in sintering of ZrO2 La2O3-B2O3 compounds are beneficial to lowering the coeffithermal expansion of ZrO2 scale,13-15 thus enhancing cient of the damage resistance of oxide scale during cyclic exposures.  this new material. But, we believe that  liquid formed during oxidation could play  an  scale,  and the  formed  IV.  Conclusions  A new approach is developed to improve the oxidation resis1500°C. A quantity  tance  of ZrC at  of  25 vol% LaB6 added to ZrC before densification. After reactively hot press1900°C,  is  ing  at  introduced LaB6 ZrB2 and a layered La-containing phase. The composite prepared by ZrC with 25 vol% LaB6 has better static oxidation resistance than ZrC at 1500°C, but shows inferior static oxi reacts with ZrC to  form  dation  resistance  compared with ZrC-40 vol% ZrB2. The enhancement of oxidation resistance of ZrC in static air after  (a)  (b)  Fig. 1.  Microstructure analysis of CBL: (a) BSE image, (b) elemental map for the distribution of La.  (a)  (b)  (c)  Fig. 2.  Macrographs of different coupons after static oxidation at 1500°C for 4 h:  (a) CBL, (b) monolithic ZrC, (c) ZrC-40 vol% ZrB2.  (a)  (b)  Fig. 3. Macrographs of CBL (a) and ZrC-40 vol% ZrB2 cyclic oxidation at 1500°C for 2 h.  (b) after  November 2011  Rapid Communications of  the American Ceramic Society  3649  \\x0c', '3650  Rapid Communications of  the American Ceramic Society  Vol. 94, No. 11  (a)  (c)  (b)  (d)  Fig. 4.  SEM micrographs of  the surface of different samples after static oxidation at 1500°C for 2 h:  (a) CBL,  (b) corresponding BSE image of  (a), (c) monolithic ZrC, (d) ZrC-40 vol% ZrB2.  1M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Cau References  sey, “Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and Zirconium Compounds,” J. Eur. Ceram. Soc., 19, 2405-14 (1999). 2A. Krajewski, L. D’Alessio, and G. De Maria, “Physico-Chemical and Thermophysical Properties of Cubic Binary Carbides,” Cryst. Res. Technol., 33, 341-74 (1998). 3H. O. Pierson, Handbook  and Nitrides. William  of Refractory Carbides  Andrew Publishing/Noyes, Westwood, NJ, 1996, 68pp. 4H. Wiedemeier and M. Singh, “Thermochemical Modelling of Interfacial Reactions in Molybdenum Disilicide Matrix Composites,” J. Mater. Sci., 27, 2974-8 (1992). 5M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, “Oxidation-Based Materials Selection for 2000°C+ Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience,” J. Mater. Sci., 39, 5887-904 (2004). 6R. F. Voitovich and E. A. Pugach, “High-Temperature Oxidation of ZrC and HfC,” Powder Metall. Met. Ceram., 12, 916-21 (1973). 7S. Shimada, M. Nishisako, M.  and K. Yamamoto,  “Formation  Inagaki,  and Microstructure of Carbon-Containing Oxide Scales by Oxidation of Single Crystals of Zirconium Carbide,” J. Am. Ceram. Soc., 78 [1] 41-8 (1995). 8L. F. He, Z. J. Lin, Y. W. Bao, M. S. Li, J. Y. Wang, and Y. C. Zhou, “Isothermal Oxidation of Bulk Zr2Al3C4 at 500 to 1000°C in Air,” J. Mater. Res., 23, 359-66 (2008). 9L. F. He, Y. W. Bao, M. S. Li, J. Y. Wang, and Y. C. Zhou, “Oxidation of Zr2[Al(Si)]4C5 and Zr3[Al(Si)]4C6 in Air,” J. Mater. Res., 23, 3339-46 (2008). 10S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, “Improved Oxidation J. Am.  Resistance of Zirconium Diboride by Tungsten Carbide Additions,” Ceram. Soc., 91 [11] 3530-5 (2008). 11S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, “Oxidation of Zirconium Diboride with Tungsten Carbide Additions,” J. Am. Ceram. Soc., 94 [4] 1198-205 (2011). 12L. Y. Zhao, D. C. Jia, Y. J. Wang, J. C. Rao, Z. H. Yang, X. M. Duan,  and Y. Zhou, “ZrC-ZrB2 Matrix Composites with Enhanced Toughness Prepared by Reactive Hot Pressing,” Scripta Mater., 63, 887-90 (2010). 13I. N. Chakraborty, J. E. Shelby, and R. A. Condrate, “Properties and Structure of Lanthanum Borate Glasses,” J. Am. Ceram. Soc., 67 [12] 782-5 (1984). 14I. N. Chakraborty and D. E. Day, “Effect of R3+ Ions on the Structure and Properties of Lanthanum Borate Glasses,” J. Am. Ceram. Soc., 68 [12] 641-5 (1985). 15I. N. Chakraborty, D. E. Day, J. C. Lapp, and J. E. Shelby, “Structurein Lanthanide Borate Glasses,” J. Am. Ceram. Soc., 68 [7]  Property Relations 368-71 (1985).  h  Fig. 5.  XRD patterns of the oxide scale on CBL after static oxidation at 1500°C for 2 h (a) and 4 h (b).  LaB6 addition is attributed to the formation of ZrB2. Microstructure analysis and cyclic oxidation test indicate that the  formed La-containing  phase  in the  composite prepared by  ZrC with 25 vol% LaB6 tive density of oxide scale during oxidation and in enhancing  to increasing the  is beneficial  rela the oxide  scale  stability during cyclic  exposures. Combining  our previous results that  the layered La-containing phase can  enhance  fracture  of LaB6 important significance for application of ZrC as UHTCs.  toughness  addition  of ZrC,  has  \\x0c']"
},{
  "_id": 98,
  "PDF": "Improved Oxidation Resistance of Zirconium Diboride by Tungsten Carbide Additions.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 91 [11] 3530 - 3535 (2008)  DOI: 10.1111/j.1551-2916.2008.02713.x  r 2008 The American Ceramic Society  Improved Oxidation Resistance of Zirconium Diboride by Tungsten  Carbide Additions  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, Missouri 65409  Shi C. Zhang,*,w  Greg E. Hilmas,* and William G. Fahrenholtz**  The effect of tungsten carbide (WC) additions on the oxidation resistance of zirconium diboride (ZrB2) was studied. Zirconium diboride ceramics, nominally pure and with the addition of  4 mol% WC, were densified by pressureless sintering. During  densification,  the WC that was added to the ceramic dissolved  a  into  solid  forming  solution. Oxidation  the ZrB2 matrix, behavior was evaluated using thermal gravimetric analysis up to 15001C and isothermal furnace oxidation at 15001 and 16001C. The WC additions reduced both the weight gain and the oxide scale thickness. After oxidation at 16001C for 3 h, the mass gain decreased from 22 mg/cm2 for nominally pure ZrB2 to 8 mg/cm2 for the WC-containing ceramic. Analysis showed that the WC additions improved the oxidation resistance of  ZrB2 by promoting densification of during oxidation through liquid-phase sintering.  the zirconia scale formed  I.  Introduction  Z IRCONIUM diboride (ZrB2) is attractive for a variety of aerospace applications. ZrB2 has a unique combination of properties, including a very high melting point (430001C), the lowest theoretical density (6.09 g/cm3) among ultra-high temperature ceramics (UHTCs), high strength, and high-elastic modulus.1,2  As a result, ZrB2-based ceramics are candidates for applications such as leading edges and thermal protection systems for reus able atmospheric re-entry vehicles, hypersonic flight vehicles, and rocket propulsion systems.3-5 These materials have also been used in other high-temperature structural applications.6,7  As with other nonoxide ceramics, ZrB2 oxidizes when exposed to air at elevated temperatures. The oxidation behavior of ZrB2 has been studied since the 1960s.8-12 In general, the oxidation  can be  of ZrB2 Reaction (1), which leads to measurable mass gain above about 8001C13  stoichiometric according to  considered to be  ZrB2 ðcrÞ þ 5=2 O2 ðgÞ ! ZrO2 ðcrÞ þ B2O3 ðl Þ  (1)  ZrB2 oxidation can be divided into three stages as the temperature increases, based on the behavior of the B2O3 that is formed.12 Below B12001C, molten B2O3 forms a continuous film, which results in parabolic mass gain kinetics because one of  the reactants (oxygen) must diffuse through the product layer (B2O3) to react with ZrB2. At intermediate temperatures (B12001 to B14001C), paralinear kinetics are observed due to partial evaporation of the B2O3. Above B14001C, evaporation of B2O3 is rapid, leaving behind a porous ZrO2 layer that does not protect from further oxidation. Recently, Parthasarathy et al.14 ZrB2  N. Jacobson—contributing editor  Manuscript No. 24439. Received March 19, 2008; approved August 14, 2008.  This work was  funded by the Air Force Research Laboratory on contract number  FA8650-04-C-5704 through the Center for Aerospace Manufacturing Technologies at the  Missouri University of Science and Technology.  *Member, The American Ceramic Society.  **Fellow, The American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: scz@mst.edu  proposed a model  to describe the oxidation behavior and final  oxidized microstructure for ZrB2, TiB2, and HfB2. The model, and previous experimental results,5,8,12 revealed that the ZrO2 oxide scale was not protective at elevated temperatures because it was porous.12 The porosity, combined with the columnar micro structure of the ZrO2 scale, offers channels for oxygen transport directly and rapidly to the reaction interface.15 Oxygen transport  is much faster  through the  resulting pores as  compared with  diffusion through the ZrO2 crystal structure. The strategy most often used to improve the oxidation resistance of ZrB2-based materials above 14001C is the addition of Si-containing compounds8,16-26 such as SiC, MoSi2, or TaSi2. In the late 1960s, Kaufman and colleagues27 reported that addi tions of 5-30 vol% SiC to ZrB2 resulted in a significant decrease (B12001 in the oxidation rate of ZrB2 at intermediate to B14001C) and elevated (above B14001C) temperatures due to  the formation of an adherent, protective layer of borosilicate  glass.18-23,27  However, the experimental results indicated that island-like, porous ZrO2 deposits grew during oxidation,23 offering a rapid transport path for oxygen in some cases. At temperatures above B16001C, or under  the flowing gas envi ronments encountered in proposed applications,  the protective  SiO2-based glass layer can be removed, leaving behind a porous ZrO2 skeleton.2 Formation of a dense ZrO2 layer during oxidation of ZrB2 materials, with or without SiC additions, may improve their oxidation resistance.  This research focuses on a method to improve the oxidation  resistance of refractory diborides without the addition of com pounds that form SiO2 when oxidized. Specifically, the goal is to improve the oxidation resistance of ZrB2 at elevated temperatures by promoting the formation of a dense ZrO2 scale on the surface of ZrB2.  II.  Experimental Procedures  (1)  Powder Preparation and Characterization  ZrB2 (Grade B, H.C. Starck, Newton, MA) was used to prepare the materials for this study. Tungsten additions were accom plished by adding tungsten carbide  (WC) powder  (CERAC,  Milwaukee, WI) to some batches. For some samples, the ZrB2 powders were attrition milled (Model 01-HD, Union Process,  Akron, OH) using WC milling media to reduce the particle size.  After milling, the average particle sizes were calculated from sur face area values determined by nitrogen adsorption (NOVA 1000,  Quantachrome, Boynton Beach, FL). ZrO2 Hill, MA) and tungsten oxide (WO3) (Aldrich Chemical, Milwaukee, WI) powders were used to study the sintering behavior of  (Alfa Aesar, Ward  the possible oxidation products. The purity, particle  size, and  suppliers of ZrB2, WC, ZrO2, and WO3 are reported in Table I.  (2)  Sample Preparation  To enhance the densification of ZrB2, 2 wt% B4C (Grade HS, H.C. Starck) was added to all of the batches.28 The as-received  2  ZrB2, wt% B4C, were dispersed in methyl ethyl ketone by ball milling for 24 h  and WC (for W-containing  batches)  with a dispersant (DISPERBYK-110, BYK-Chemie Co., Wesel,  Germany). Next, 1 wt% binder (Qpac-40, Empower Materials,  3530  \\x0c', 'Newark, DE) was added and the mixture was milled for another  24 h. After ball milling, the slurry was dried and ground to form  granules for dry pressing. Cylindrical disks 19 mm in diameter  were formed by uniaxial pressing at 18.6 MPa (2.7 ksi), followed  by cold isostatic pressing at 315 MPa (45 ksi).  Compacted pellets were densified by pressureless sintering. Pellets were heated at 101C/min in a graphite crucible to temperatures ranging from 18501 to 20501C using a graphite element  furnace (3060-FP20, Thermal Technology, Santa Rosa, CA). The furnace atmosphere was a mild vacuum (B20 Pa) for temperatures of 16501C or below, but was switched to flowing argon (B105 Pa) for temperatures above 16501C. Hold times were from  2 to 3 h for  sintering. The phases present after heating were  analyzed by X-ray diffraction (XRD; XDS 2000, Scintag Inc.,  Cupertino, CA). The bulk densities of sintered specimens were  determined using  the Archimedes method. Relative densities  were calculated by dividing the measured bulk densities by the  theoretical densities, which were calculated for each composition  based on the amounts of B4C and WC in the batch compositions. Densified pellets with a diameter of B18 mm and a thickness of B5 mm were produced. Specimens with dimensions of 5 mm \\x02 4 mm \\x02 3 mm were diced from the pellets for oxidation tests.  (3)  Oxidation Testing  The mass gain was characterized as a function of temperature for  ZrB2-based ceramics with and without WC additions by thermal gravimetric analysis (TGA; Netzsch Sta 409C/CD, Selb, Ger many). Mass gain was also characterized as a function of time at 15001 or 16001C using static testing in a MoSi2 resistance heated horizontal mullite tube furnace (Model 0000543 Rapid Temper ature Furnace, CM Inc., Bloomfield, NJ) equipped with gas-tight  end caps. Before oxidation, specimens were cleaned in an ultra sonic bath with acetone. For TGA testing, the rectangular spec imens were placed in an Al2O3 crucible that had been notched in two places near the bottom of the crucible to allow air to flow  freely around the specimens. For isothermal oxidation studies,  similar  rectangular  specimens were placed on a zirconia setter  with ridges to minimize the contact area between the specimens and the setter. Specimens were heated at B51C/min to 15001 or 16001C and held for 1, 2, or 3 h in a flowing air atmosphere.  After  cooling  to room temperature,  the mass was measured (mg/cm2) were calculated  again. The normalized weight gains  from mass gain and calculated surface area of the specimens.  (4)  Sample Characterization  The thicknesses of the resulting oxidation layers were measured  from polished cross sections. The microstructures of the oxide scale  of some oxidized specimens were observed using scanning elec tron microscopy (SEM; S-570, Hitachi, Tokyo, Japan), and el emental distribution maps were  collected in the SEM using  energy-dispersive  spectroscopy  (EDS; E2V,  Scientific  Instru ments, Buckinghamshire, U.K.).  III.  Results and Discussion  (1)  Oxidation of Nominally Pure ZrB2  The TGA results for nominally pure ZrB2 materials (498% of theoretical density) indicated a weight gain of 6.7 mg/cm2 when  the temperature reached 15001C (Fig. 1). In addition, the mass  gain increased more 15001C,  rapidly  as  the  temperature  approached  indicating  that  the  oxidation  rate  increased  as  the  temperature increased. Both the Parthasarathy model, and the  previous experimental studies, have concluded that oxidation of  nominally pure ZrB2 in air at an elevated temperature produced a porous scale. Figure 2 shows a typical cross-section micro structure of the ZrB2 prepared from as-received ZrB2 powder oxidized at 12001C for 2 h in air. The porous scale consists of  columnar ZrO2 grains extending to the surface. The channels in the ZrO2 outer scale offer a rapid path for oxygen transport to the interface between ZrO2 and the underlying, unoxidized increasing the density of  ZrB2. Hence, the ZrO2 scale may improve the oxidation resistance of ZrB2 in this temperature range.  (2)  Liquid-Phase Sintering  Liquid-phase sintering is one approach that may increase the  relative density of the ZrO2 scale that forms during oxidation of ZrB2 at elevated temperatures. Phase equilibrium diagrams29 show that ZrO2 forms a liquid phase with WO3 at temperatures above 12601C. The minimum amount of WO3 needed to form a liquid phase with ZrO2 at 12601C is B4 mol%. To confirm that WO3 can promote densification of ZrO2 by liquid-phase sintering, a series of experiments were conducted using mixtures of  Table I.  Raw Material Characteristics  Materials  Grade  (purity%)  Particle  size (mm)  Surface  area  (m2/g)  Oxygen  content  (%)  Supplier  ZrB2 B4C WC  B499  2  1  0.9  H.C. Starck  HS 499.5  0.8  15.8  1.3  H.C. Starck  o1  CERAC  ZrO2 WO3  991  0.3  3.2  —  Alfa  99.995  —  Aldrich  Fig. 1.  Thermal gravimetric analysis results for nominally pure zirco nium diboride (ZrB2) and ZrB2 containing 4 mol% tungsten carbide (WC) heated at 101C/min to 15001C in flowing air.  Fig. 2. Fracture surface of the scale on nominally pure zirconium diboride (ZrB2) with 2 wt% B4C that was oxidized at 12001C in air for 2 h.  November 2008  Oxidation Resistance of Zirconium Diboride  3531  \\x0c', '3532  Journal of the American Ceramic Society—Zhang et al.  Vol. 91, No. 11  Fig. 3.  Fracture surface of ZrO2 containing 4 mol% tungsten oxide heated to 14501C showing evidence of tungsten-rich phase.  liquid-phase sintering of  the ZrO2 by a  ZrO2 and WO3. Pellets of ZrO2 (B0.3 mm starting particle size) with WO3 additions of 4, 6, and 10 mol% were sintered at 14501C for 3 h in air. All of the pellets reached full density. For  ZrO2 containing 4 mol% WO3, SEM showed that the ZrO2 particles grew and developed a round morphology. The ZrO2 particles, which were initially B0.3 mm in diameter, grew to B20 mm in diameter (Fig. 3) at 14501C, roughly 70 times their initial  size. Chemical analysis by EDS confirmed that a W-rich phase  surrounded the round ZrO2 grains. Microstructure and chemical analyses indicated that a WO3 addition of 4 mol% led to liquidphase formation in ZrO2. Further, the liquid phase appeared to promote densification of the mixture by liquid-phase sintering.  The volatilization of WO3 was also considered. If, like B2O3, the WO3 formed during oxidation evaporated rapidly, then it may not be an effective liquid-phase former for sintering ZrO2. The vapor pressures of various WO3 species (e.g., WO3, (WO3)2, etc.) were calculated using thermodynamic software30 and the  results are summarized in Table II. From the calculation, (WO3)3 had the highest vapor pressure at 15001C, which was 112 Pa (1.1 \\x02 10\\x003 atm.). Based on observations of the sintered mixture of ZrO2 and 4 mol% WO3 shown in Fig. 3, the W-rich oxide phase did not appear to undergo significant evaporation  despite the relatively high vapor pressure. This may be due to  the  low W concentration in the  system, which would reduce  the activity of W in the ZrB2 or the relatively low fraction of the outer surface composed of WO3, which would reduce the evaporation rate. Thus, the actual vapor pressure of the (WO3)n species may be lower than the values shown in Table II.  (3)  Solid Solution Formation  Previous research has shown that oxide impurities inhibit dens ification of ZrB2 by promoting coarsening of ZrB2 at temperatures below which it can be densified by solid-state sintering.28  Therefore, the addition of WO3 directly to the ZrB2 would likely inhibit densification and perhaps have other adverse effects on  the properties of the resulting ceramics. Therefore, it is desirable  Table II.  Vapor Pressure of Tungsten Oxide (WO3)n Species in Equilibrium with WO3 Solid  Vapor species  WO3 (WO3)2 (WO3)3 (WO3)4  P at 15001C (Pa)  \\x003  1.08 \\x02 10 12.2  112  35.3  P at 16001C (Pa)  \\x003  8.25 \\x02 10 75.2  565  20.4  to add the W source in a form that can be uniformly distributed  inhibiting the densifi in the ZrB2 before sintering and without cation of ZrB2. In addition, the precursor should produce WO3 during ZrB2 oxidation. In a previous study of pressureless sintering of ZrB2,28 WC, present as an impurity from particle size reduction using co-bonded WC grinding media, dissolved into lattice.28,31-33 Figures 4(a) and (b)  the ZrB2 patterns of ZrB2 containing 4 mol% WC before and after pressureless sintering at 20501C for 3 h in Ar. The WC phase  show the XRD  that was detected in the green pellets before sintering (Fig. 4(a))  disappeared after sintering and only ZrB2 was observed by XRD (Fig. 4(b)). The peaks in Fig. 4(b) can be indexed to the XRD  powder diffraction file  card (PDF card number 34-0423)  for  2y  the peaks  hexagonal ZrB2. However, higher values of compared with pure ZrB2. Because W (1.4 A˚ ) has a smaller covalent radius than zirconium (Zr) (1.6 A˚ ),34 substitution of W into the ZrB2 lattice is expected to decrease the average unit-cell size, and shift the diffraction peaks to higher 2y values in Fig. 4.  shifted to slightly  are  Thermodynamic analysis indicated that  the standard Gibbs  free  energy of  reaction for  the oxidation of W to WO3 was favorable over the entire temperature range studied, from room to 20001C. XRD phase analysis of an oxidized  temperature  mixture of 10 mol% WC and 90 mol% ZrB2 powders showed that monoclinic ZrO2 and hexagonal WO3 were produced during oxidation at 15001C in air. Because the WC and ZrB2 form a solid solution, it is reasonable to predict that an inti mately mixed oxide scale containing both ZrO2 and WO3 would form during oxidation. Therefore, WC seems to be an accept able precursor for WO3 because WC forms a solid solution with ZrB2 and it oxidizes to form WO3 when ZrB2 is oxidized. As discussed above, the condensed WO3 appeared to be stable over the temperature range reported in this study. One final consid eration was the volume change during oxidation. When W is oxidized to WO3, the volume increases by B230%, which means the oxidation of a small amount of W produces a large volume  that  form protective oxide scales,  of WO3. Materials Al, often show a substantial volume increase upon oxidation. increases by B25% when it oxidizes For Al, the volume form a-Al2O3. Tailoring W-containing ZrB2 is oxidized may also be beneficial formation of a dense ZrO2-based protective scale.  increase when  such as  volume  overall  the  to  for  the  (4)  Oxidation Behavior  TGA results  for ZrB2-containing 4 mol% WC are compared with the oxidation of nominally pure ZrB2 in Fig. 1. Analysis indicates that the weight gain for the ZrB2 ceramic with 4 mol% WC was 4.4 mg/cm2 compared with 6.7 mg/cm2 for nominally  \\x0c', 'November 2008  Oxidation Resistance of Zirconium Diboride  3533  Fig. 5. Weight gain for nominally pure zirconium diboride (ZrB2) and tungsten-containing ZrB2 after isothermal oxidation at 15001 and 16001C for 60, 120, or 180 min.  pure ZrB2 ceramics increased from B50 mm after a 1-h hold to B500 mm after 3 h. Under the same oxidation conditions, the  thickness of the ZrO2 scale for the WC-containing ZrB2 was only B25 mm for 1 h, with an increase to B100 mm after 3 h.  Hence,  the WC-containing ceramics exhibited better oxidation  resistance than nominally pure ZrB2 based on both mass gain and scale thickness. The reduced weight gain can be attributed  to the higher density of  the ZrO2 outer  scale at  the higher  temperature. Remarkably, at 16001C,  the weight gains for ZrB2 ceramics containing 4 mol% WC were substantially lower for the same isothermal hold times as compared with the results at 15001C (Fig. 5). For example, after 3 h at 16001C, the WC-containing ZrB2 gained only 8 mg/cm2, compared with 11 mg/cm2 for the same W-containing ZrB2 oxidized at 15001C for 3 h. Based on this analysis, WC is more effective at improving the oxidation  resistance of ZrB2 at higher temperatures, probably due to more effective liquid-phase sintering of the ZrO2 scale at 16001C.  (5)  Microstructure Characterization  To confirm the beneficial role of the W additions, cross sections  of the ZrO2 scales of nominally pure ZrB2 and the ZrB2 containing WC that were oxidized at 16001C in air for 2 h were  characterized using SEM and compared (Fig. 6). For the nominally pure ZrB2, the ZrO2 oxidation scale was B1.15 mm thick (Fig. 6(a)) and consisted of columnar ZrO2 grains with open pore channels between them (Fig. 6(b)). The open channels offer  a path for  rapid oxygen transport  to the ZrO2/ZrB2 where oxidation occurs. In contrast, the ZrO2 oxidation scale on the WC-containing ZrB2 was o0.3 mm thick (Fig. 6(c)) with a dense microstructure that consisted of equiaxed ZrO2 grains (Fig. 6(d)). This microstructure appears to be a barrier  interface  to oxygen transport. Therefore, the weight gain after oxidation at 16001C for 2 h in air was lower, only B7 mg/cm2, compared  with the nominally pure ZrB2, which had a weight gain of B14 mg/cm2 under the same conditions. W and Zr elemental  maps of  the area shown in Figs. 6(e) and (f)  confirmed the  presence of W in the dense surface layer. Thus, the microstruc Table III.  Comparison of the Thicknesses of the Outer  Scale Produced by Oxidation of Nominally Pure Zirconium  Diboride (ZrB2) and ZrB2-Containing 4 mol% Tungsten Carbide (WC) at 15001C  Materials  Nominally pure ZrB2 ZrB214 mol% WC  1 h  50 mm 25 mm  2 h  3 h  150 mm 50 mm  500 mm 100 mm  Fig. 4.  zirconium diboride (ZrB2) taining 4 mol% tungsten carbide (WC) (a) before and (b) after sintering  X-Ray diffraction pattern of  con and (c) after oxidation.  pure ZrB2, a reduction of about 35% in the mass gain. However, the TGA curves for nominally pure ZrB2 and WC-containing ZrB2 both showed the same trend of increasing mass gain as the temperature approached 15001C. The full benefit of WC addi tions is not apparent in this experiment in which the specimens  were heated and cooled with no isothermal  treatment. In this  case,  lower weight  the WC-containing ZrB2 was attributed to the formation of a stable oxide that filled the voids  gains  for  between the ZrO2 grains. The volume increase associated with W oxidation should produce an oxide that fills the voids  between the ZrO2 grains, much as B2O3 would provide oxidation protection at lower temperatures.  To fully understand the benefits of WC additions to ZrB2, isothermal oxidation studies were conducted at 15001 and 16001C. Figure 5 compares the weight gains of nominally pure ZrB2 and ZrB2 containing 4 mol% WC at 15001 and 16001C for hold times up to 3 h. The WC-containing ZrB2 ceramics exhibited a lower weight gain at both temperatures for all hold times  compared with the nominally pure ZrB2. For example, after 3 h at 15001C, the WC-containing ZrB2 gained 11 mg/cm2 compared with 15 mg/cm2 for the nominally pure ZrB2. Cross isothermally oxidized at 15001C  both materials  sections  of  were observed using SEM. The thicknesses of the oxidized outer  scales were measured from the SEM images  and compared  in Table  III. The  thickness of  the ZrO2  scale  for nominally  \\x0c', '3534  Journal of the American Ceramic Society—Zhang et al.  Vol. 91, No. 11  (a)  (c)  (b)  (d)  (e)  (f)  Fig. 6.  Scanning electron micrographs of ZrO2 outer scale for nominally pure zirconium diboride (ZrB2) (a, b) and tungsten carbide-containing ZrB2 (c, d) and energy-dispersive spectroscopy compositional maps (e, f) corresponding to the areas in (c) and (d).  ture and composition of  the oxide were  consistent with the  sidering all of  these factors,  it  is not  surprising that  the WO3  formation of a W-containing oxide phase, which apparently  phase was not detected using XRD.  promoted liquid-phase densification of the ZrO2 scale. The phase composition of the outer scale was also investigated further using XRD analysis. After oxidation at 15001C for  3 h, only monoclinic ZrO2 was detected on the surface of ZrB2 containing 4 mol% WC (Fig. 4(b)). However, additional SEM  analysis, including W element mapping by EDS, of a specimen oxidized at 16001C (Fig. 6e), indicated that some W was present  in the ZrO2 outer scale. Because crystalline WO3 was not observed by XRD, its detection by SEM/EDS can be attributed to (1) according to the phase diagram,29 the solu several factors:  bility of WO3 in ZrO2 is about 2 mol% at room temperature; (2) if no WO3 were lost during oxidation testing, only B2 mol% WO3 would remain and so the relative intensities of the WO3 peaks would be quite low compared with ZrO2, perhaps below the XRD detection limit; (3) the major peaks of monoclinic  WO3, which are from the (002) and (110) planes, are located at 2y values of 23.2511 and 24.2251, very close to the (110) and (011) planes of monoclinic ZrO2, which has peaks at 24.0681 and 24.4611; and (4) some of the WO3 may have evaporated during to o2 mol%. Conoxidation, reducing the final WO3 content  IV.  Conclusion  A new approach was developed to improve the oxidation resis tance of ZrB2. A W compound, WC in this to ZrB2 before densification by pressureless oxidation, the presence of WO3 in the oxide liquid-phase sintering of the ZrO2 increased its density, and decreased the rate of  scale, which modified  sintering. During  case, was added  scale  resulted in  its  microstructure,  oxygen transport. As a result, WC-containing ZrB2 had improved oxidation resistance compared to nominally pure ZrB2 as measured by mass gain and oxide scale thickness. The theoretical  considerations  and  experimental  results  point  to  three main  criteria that could be used to select other additives  that might  similarly enhance the oxidation resistance of Zr-based UHTCs:  (1) WC formed a solid solution with ZrB2 and was uniformly distributed in the sintered ZrB2 ceramic. (2) During oxidation, WC-containing ZrB2 simultaneously oxidized to form both WO3 and ZrO2. WO3 formed a liquid  \\x0c', 'phase with ZrO2, which enhanced the densification of ZrO2 to form a dense outer scale with an equiaxed microstructure. In  contrast,  the ZrO2 formed on nominally pure ZrB2 was porous and consisted of columnar grains.  scale that  (3)  The volume increase (230%) associated with oxidation  of W to form WO3 enhanced the increase in total volume from 13% for formation of ZrO2 during oxidation of nominally pure ZrB2 to 23% for formation of WO3 and ZrO2 during oxidation of ZrB2 containing 4 mol% WC. (4) The criteria for selecting other additives that may pro duce similar  improvements in oxidation resistance of diboride  ceramics are as follows: (a) forms a liquid phase with ZrO2 at moderate temperatures (below 14001C) and at low amounts; (b)  dissolves into the diboride matrix; and (c) has a large volume  increase upon conversion of the metal to the oxide.  References  1W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. 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Tanaka,  ‘‘Development of ZrB2-Graphite Protective Sleeve for Submerged Nozzle,’’ Taikabutsu Overseas, 14 [2] 39-43 (1992). 7O. Kida and Y. Segawa,  ‘‘ZrB2 Composite Sintered Material’’; U.S. Patent  4,636,481, Jan 13 1987. 8P. T. B. Shaffer,  ‘‘An Oxidation Resistant Boride Composition,’’ Am. Ceram.  Bull., 41 [2] 96-9 (1962). 9A. K. Kuriakose and J. L. Margrave,  ‘‘The Oxidation Kinetics of Zirconium  Diboride and Zirconium Carbide at High Temperatures,’’ J. Electrochem. Soc.,  111 [7] 827-31 (1964). 10J. B. Berkowitz-Mattuck,  ‘‘High Temperature Oxidation. III. Zirconium and  Hafnium Diborides,’’ J. Electrochem. Soc., 113 [9] 908-14 (1966). 11L. Kaufman, E. V. Clougherty, and J. B. Berkowitz-Mattuck,  ‘‘Oxidation  Characteristics of Hafnium and Zirconium Diboride,’’ Trans. Metall. Soc. AIME,  239 [4] 458-66 (1967). 12W. G. Fahrenholtz,  ‘‘The ZrB2 Volatility Diagram,’’ J. Am. Ceram. Soc., 88  [12] 3509-12 (2005).  13A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Evolution of Structure  During the Oxidation of Zirconium Diboride-Silicon Carbide 15001C,’’ J. Eur. Ceram. Soc., 27 [6] 2495-501 (2007). 14T. A. Parthasarathy, R. A. Rapp, M. Opeka, and R. J. Kerans, ‘‘A Model for  in Air  up  to  the Oxidation  of ZrB2, HfB2,  and TiB2,’’ Acta Mater.,  55  [17]  5999-6010  (2007). 15J. Li and T. J. Lenosky,  ‘‘Thermochemical and Mechanical Stabilities of the  Oxide Scale of ZrB21SiC and Oxygen Transport Mechanisms,’’ J. Am. Ceram. Soc., 91 [5] 1475-80 (2008). 16E. V. Clougherty, R. L. Pober, and L. Kaufman,  ‘‘Synthesis of Oxidation  Resistant Metal Diboride Composites,’’ Trans. Metall. Soc. AIME, 242 [6] 1077-  82 (1968). 17W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of an SiC Addition on  the Oxidation of ZrB2,’’ Am. Ceram. Bull., 52 [8] 612-6 (1973). 18A. Bongiorno, C. J. Fo¨ rst, R. K. Kalia, J. Li, J. Marschall, A. Nakano, M. M.  Opeka, I. G. Talmy, P. Vashishta, and S. Yip,  ‘‘A Perspective on Modeling Ma terials in Extreme Environments: Oxidation of Ultrahigh-Temperature Ceramics,’’  MRS Bull., 31 [5] 410-8 (2006). 19W. G. Fahrenholtz,  ‘‘Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 20M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurfaces: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 21F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 22C. R. Wang, J. M. Yang, and W. Hoffman, ‘‘Thermal Stability of Refractory  in Dry  Carbide/Boride Composites,’’ Mater. Chem. Phys., 74 [3] 272-81 (2002). 23S. N. Karlsdottir and J. W. Halloran,  ‘‘Rapid Oxidation Characterization of  Ultra-High Temperature Ceramics,’’ J. Am. Ceram. Soc., 90 [10] 3233-8 (2007). 24D. Sciti, M. Brach, and A. Bellosi,  ‘‘Long-term Oxidation Behavior and Me chanical Strength Degradation of a Pressureless Sintered ZrB2-MoSi2 Ceramic,’’ Scr. Mater., 53, 1297-302 (2005). 25E. Opila, S. Levine, and J. Lorincz,  ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39 [19]  5969-77 (2004). 26I. G. Talmy, J. A. Zaykoski, M. M. Opeka, and A. H. Smith,  ‘‘Properties of  Ceramics in the System ZrB2-Ta5Si3,’’ J. Mater. Res., 21 [10] 2593-9 (2006). 27E. V. Clougherty, R. L. Pober, and L. Kaufman, ‘‘Synthesis of Oxidation  Reisistance Metal Diboride Composites,’’ Trans. TMS-AIME, 242 [6] 1077 (1968). 28S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Pressureless Densificat ion of Zirconium Diboride with Boron Carbide Additions,’’ J. Am. Ceram. Soc.,  89 [5] 1544-50 (2006). 29L. L. Y. Chang, M. G. Scroger, and B. Phillips,  ‘‘Diagram Zr-213’’; p. 148 in  Phase Diagrams for Zirconium and Zirconia Systems, Edited by H. M. Ondik, and  H. F. McMurdie. The American Ceramic Society, Westerville, OH, 1998. 30HSC Chemical 5.1, ESM Software. Hamilton, OH 2002. 31B. Post, F. W. Glaser, and D. Moskowitz,  ‘‘Transition Metal Diborides,’’  Acta Metall., 2 [1] 20-5 (1954). 32W. A. Zdaniewski,  ‘‘Solid  Solubility Effect  on  Properties  of Titanium  Diboride,’’ J. Am. Ceram. Soc., 70 [11] 793-7 (1987). 33A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Pressureless  Sintering of Zirconium Diboride,’’ J. Am. Ceram. Soc., 89 [2] 450-6 (2006). 34F. Laves,  ‘‘Theory of Alloy Phases’’; pp. 124-98 American Society of Metals,  Cleveland, OH, 1956.  &  November 2008  Oxidation Resistance of Zirconium Diboride  3535  \\x0c']"
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  "_id": 99,
  "PDF": "Improved processing and oxidation-resistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids.pdf",
  "Text": "['Materials Science and Engineering A 464 (2007) 216-224  Improved processing and oxidation-resistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids  Sung S. Hwang, Alexander L. Vasiliev, Nitin P. Padture  ∗  Department of Materials Science and Engineering, The Ohio State University, Columbus, OH 43210, USA  Received 21 September 2006; received in revised form 28 January 2007; accepted 1 March 2007  Abstract  We have studied the hot-pressing behavior of ZrB2 /SiC ultra-high temperature ceramics (UHTCs) as a function of: (i) SiC starting-powder size, (ii) SiC vol%, (iii) ZrO2 doping, and (iv) colloidal dispersion of ZrB2 /SiC powder mixtures. It has been found that the addition of SiC promotes densification of ZrB2 at a moderate hot-pressing temperature of 1650 C. It has also been found that ball-milling of the ZrB2 /SiC starting-powder mixtures using ZrO2 balls media results in the doping of the powder mixture with ZrO2 , which promotes hot-pressing densification. Reduction in the SiC starting-powder size, and colloidal dispersion of the powders, both have been found to promote hot-pressing densification of ZrB2 /SiC materials; the highest density achieved in such ZrB2 /SiC ceramics is 99.9%. Detailed microstructural characterization of the ZrB2 /SiC ceramics using electron microscopy shows that some of these materials contain a Zr(O,B)2 phase, and amorphous films at interphase interfaces. Oxidation studies reveal that SiC grain-size reduction results in improved oxidation-resistance in ZrB2 /SiC materials. The ZrB2 /SiC ceramics produced here possess modest hardness and toughness properties. The results presented here point to a new strategy for improving processing and oxidation-resistance of ZrB2 /SiC materials: dispersion and reduction of SiC grains. © 2007 Elsevier B.V. All rights reserved.     Keywords: Ceramics; Composites; Oxidation; Processing; Microstructure  1.  Introduction  Zirconium diboride (ZrB2 ) is a leading candidate material for use in critical external surfaces of future aerospace re-entry crafts, such as hypersonic aircraft and reusable launch vehicles [1,2]. This is primarily because ZrB2 , a so-called ultra-high temperature ceramic (UHTC), has an excellent combination of mechanical, physical, thermal-shock, and oxidation-resistance properties. It has been shown that addition of SiC dispersoids to ZrB2 ceramics results in significant improvements in oxidationresistance and mechanical properties compared to ZrB2 alone [1,3-9]. ZrB2 /SiC materials are generally fabricated using hotpressing at high temperatures, in the range 1900 C [7,8] to 2100 C [3]. Hot-pressing at lower temperatures (1600-1750 C) has been shown to result in dense ceramics, but with additions of TaSi2 [6] or Si3N4 [5]. The average particle sizes of the ZrB2 and SiC starting powders used are between           ∗  Corresponding author. Tel.: +1 614 247 8114; fax: +1 614 292 1537.  E-mail address: padture.1@osu.edu (N.P. Padture).  0921-5093/$ - see front matter © 2007 Elsevier B.V. All rights reserved.  doi:10.1016/j.msea.2007.03.002  1 and 5 \\u242em, and the SiC content varies from 5 to 25 vol% SiC. Although these processing conditions result in fully dense ZrB2 /SiC materials, systematic studies on the effects of SiC starting-powder reduction and ZrO2 doping on the hot-pressing behavior of ZrB2 /SiC materials are lacking. Also lacking is detailed characterization of microstructures of ZrB2 /SiC materials. Recent oxidation studies have shown that the simultaneous oxidation of ZrB2 and SiC in ZrB2 /SiC materials results in a less volatile silica-rich surface scale, in place of the more volatile boria surface scale in the case of pure ZrB2 [1,3,4,6,7,9]. The silica-rich surface scale is also more refractory and resistant to oxygen diffusion, making ZrB2 /SiC materials more oxidation resistant than pure ZrB2 . Furthermore, the silica-rich scale is anchored by the oxidized sub-layer consisting of an interpenetrating composite of ZrO2 /silica-rich glass. Although effects of dopants on the oxidation behavior of ZrB2 /SiC materials have been studied, microstructural effects on the oxidation behavior of ZrB2 /SiC materials have not been studied systematically. In this work, we have studied the hot-pressing behavior of ZrB2 /SiC materials as a function of: (i) SiC starting-powder  \\x0c', 'S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  217  (15 g individual batches) were respective powder mixtures ball-milled in methanol using Y2O3 -stabilized tetragonal ZrO2 (YSZ) balls (10 mm diameter) media (Tosoh Corp., Tokyo, Japan) for 24 h. All ball-milling was performed in polyethylene bottles, with powder mixture:media:methanol volume ratio of 1:2:3.2. In an experiment to study effect of powder dispersion (material Z/S-6), the 40-nm SiC powder was first dispersed in an aqueous solution of pH 12, with 2 h of ultrasonication. The pH was adjusted using NH4OH (Mallinckrodt Baker, Phillipsburg, NJ). The ZrB2 powder was then added to the dispersion. The mixture was then ball-milled for 24 h using YSZ balls media. The pH of 12 was chosen based on results from dispersion studies (not shown here), where best dispersion was observed. During a typical ball-milling run, it was found that the YSZ balls lost about 0.3 g of weight, which is assumed to be incorporated in the powder mixture. In an effort to isolate the effect of ball-milling, the powder mixture for material Z/S-7 was mixed in methanol, but without the YSZ balls media. In an additional experiment, powder mixture for material Z/S-8 was mixed in the YSZ balls media, but 0.3 g (2 wt%) of ethanol without ZrO2 powder (Alfa Aesar, Ward Hill, MA) was added to the powder mixture. All powder mixtures were dried on hot-plates while being stirred, and the resulting dried powder mixtures were crushed. Individual batches of powder mixtures (7.5 g each) were placed in BN-coated graphite dies (25.4 mm diameter), and hot-pressed (GCA Vacuum Industries, Somerville, MA) under vacuum at 1650 C for 2 h at an applied pressure of 60 MPa. All surfaces of the hot-pressed specimens were ground and cleaned. The densities of the consolidated specimens were measured using the Archimedes principle, with distilled water as the immersion medium. The measured densities are reported in Table 1, where the theoretical densities of the composites were calculated using rule-of-mixture and following density values −3 and SiC 3.217 g cm −3 . of the pure phases: ZrB2 6.085 g cm     2.2. Oxidation  Cubes 2 mm × 2 mm × 2 mm were cut out from the hotpressed materials, and all surfaces were diamond-polished to a 1 \\u242em finish using routine metallographic methods. Oxidation tests on these cubes were performed in a box furnace (Thermolyne, Dubuque, IA) at 1500 C for 10 min in air.     Fig. 1. Bright-field TEM micrograph of the 40-nm SiC starting powder.  size, (ii) SiC vol%, (iii) ZrO2 doping, and (iv) colloidal dispersion of ZrB2 /SiC powder mixtures. We have also characterized in detail the microstructures of the resulting ZrB2 /SiC materials using electron microscopy, in an attempt to elucidate densification mechanisms. Finally, microstructural effects on the oxidation behavior and mechanical properties of the resulting ZrB2 /SiC materials have been studied, in an effort to understand structure-property relations.  2. Experimental procedure  2.1. Processing  The ZrB2 starting powder used in this study was obtained from a commercial source (Grade B, H.C. Starck Corp., Newton, MA), with an average particle size of 2 \\u242em. Three types of commercial SiC starting powders with different average particles sizes were used: (i) 1.7 \\u242em (Grade B-hp SiC, H.C. Starck Corp., Newton, MA), (ii) 0.6 \\u242em (UF-15 SiC, H.C. Starck, Newton, MA), and (iii) 40 nm (experimental grade CVD SiC, Sumitomo Chemical Company, Tokyo, Japan). Fig. 1 shows a transmission electron micrograph (TEM) of the 40-nm SiC powder. Table 1 summarizes the nine powder batches that were prepared. In the case of materials Z-1, and Z/S-1 to Z/S-5 the  Table 1  Nomenclature, processing details, and microstructural parameters of ZrB2 /SiC materials hot-pressed at 1650     C under 60 MPa pressure for 2 h  Material #  SiC vol%  SiC powder size  Relative density (%)  ZrB2 grain size (\\u242em)  SiC grain size  Remarks  Z-1  Z/S-1  Z/S-2  Z/S-3  Z/S-4  Z/S-5  Z/S-6  Z/S-7  Z/S-8  0  5.7  11.4  22.4  22.4  22.4  22.4  22.4  22.4  - 1.7 \\u242em 1.7 \\u242em 1.7 \\u242em 0.6 \\u242em  40 nm  40 nm 1.7 \\u242em 1.7 \\u242em  71.6  81.7  91.5  97.9  98.6  99.6  99.9  90.9  96.4  3.4  3.6  3.6  3.9  3.7  3.4  3.4  3.4  3.6  - 1.8 \\u242em 1.9 \\u242em 1.8 \\u242em 0.8 \\u242em  80 nm  80 nm 1.8 \\u242em 1.9 \\u242em  Ball-milled  Ball-milled  Ball-milled  Ball-milled  Ball-milled  Ball-milled  Colloidal + ball-milled  Not ball-milled  ZrO2 + not ball-milled  \\x0c', 'Average silica layer thickness I* (\\u242em)  Average depleted-layer thickness II* (\\u242em)  13  4  3  2  32  9  8  7  polishing, dimpling, and ion-beam milling (DuoMill, Gatan, Pleasanton, CA). The resulting specimens were observed in a TEM (Tecnai F20, FEI, Hillsboro, OR), equipped with a XCs  1.3 mm), an atmospheric-thin-window EDS (Phoenix Systwin objective lens pole-piece (spherical aberration coefficient tem, EDAX, Mahwah, NJ), and an imaging filter (GIF 2000,  218  Table 2  S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  Hardness, toughness, and oxidation properties of dense ZrB2 /22.4 vol% SiC materials  Material #  SiC grain size  Hardness, H (GPa)  Toughness, KIC (MPa m0.5 )  Z/S-3  Z/S-4  Z/S-5  Z/S-6  1.8 \\u242em 80 nm 0.8 \\u242em 80 nm (colloidal)  15.2  16.7  19.9  21.3  * See Fig. 10 for explanation.  2.3. Characterization  3.8  3.8  3.1  2.4  Cross-sections of the hot-pressed materials and oxidized specimens were diamond-polished to 1 \\u242em finish using routine metallographic methods. Polished cross-sections of only the hot-pressed specimens were then etched using a HF:HNO3 :H2O::1:1:1 (by volume) solution for 5 s at room temperature. All cross-sections were then observed in a scanning electron microscope (SEM) (XL 30 ESEM-FEG, Philips, Eindhoven, The Netherlands). Energy dispersive spectroscopy (EDS) (EDAX, Mahwah, NJ) in the SEM was used to obtain compositional maps. In the case of the oxidized specimens, the thicknesses of the surface oxidation layers on the cross-section SEM micrographs were measured. The average values from at least 10 measurements from each material are reported in Table 2. The ZrB2 and SiC grains sizes in the hot-pressed materials were estimated from the SEM micrographs using an imageanalysis software (Clemex Vision, Clemex Technology Inc., Longueil, Canada). Approximately 100 ZrB2 and 60 SiC grains per material were used. The hot-pressed materials were analyzed by X-ray diffraction (XRD) using Cu K␣ radiation (XDS 2000, Scintag Inc., Sunnyvale, CA), to confirm the phases present. The 40-nm SiC powder was dispersed on a holey carbon grid for TEM observation. TEM samples were also prepared out of hot-pressed materials Z/S-3, Z/S-6, and Z/S-7, using routine methods involving successive steps of grinding,  Fig. 2. XRD pattern of material Z/S-3 showing the presence of ZrB2 and SiC. Joint Commission on Powder Diffraction Standards (JCPDS) powder diffraction  file numbers 75-1050 and 73-1665 were used to identify the hexagonal ZrB2 and the cubic SiC phases, respectively.  powder size. In (A), %relative densities of material Z/S-7 made from non-ball milled powder mixture and material Z/S-8 made from non-ball-milled powder  mixture but with ZrO2 additions are also shown. In (B), %relative density of material Z/S-6 made from colloidally dispersed, ball-milled powder mixture is  shown.  Fig. 3. Plots of relative density (%) of ZrB2 /SiC materials hot-pressed at 1650 under identical conditions as a function of: (A) vol% SiC and (B) SiC starting C     \\x0c', 'S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  219  Gatan, Pleasanton, CA) with electron energy loss spectrometer (EELS). The TEM was operated at an accelerating voltage of 200 kV.  indentation cracks with the microstructure were also examined in the SEM.  2.4. Mechanical testing  Polished cross-sections of only the dense, hot-pressed materials (Z/S-3, Z/S-4, Z/S-5, and Z/S-6) were indented using a Vickers diamond pyramid at a contact load of 1 kg (9.8 N) (Micromet II, Buehler, Lake Bluff, IL). The indentation sites were examined using an optical microscope, and the indentation impression diagonals were measured for hardness (H) calculations. The same specimens were also Vickers-indented using a higher load (P) of 10 kg (98 N) (Zwick, Ulm, Germany) for toughness measurements. Radial cracks emanating from the Vickers indentation sites (2c) were measured in the optical microscope. The indentation toughness values of the specimens were calculated using the following relation [10]: −1.5 , where E was estimated using the Knoop indentation method (indentation load 9.8 N) [11]. The average H and KIC toughness values, from at least five indentation per material, are reported in Table 2. Interactions of the  KIC = 0.016(E/H)0.5Pc  3. Results  3.1. Microstructures  Fig. 2 is an example of a typical XRD pattern, in this case for material Z/S-3, showing the presence of only two phases, ZrB2 (hexagonal) and ␤-SiC (cubic). A comparison of density data in Table 1 and Fig. 3A for materials Z-1, Z/S-1, Z/S-2, Z/S-3 clearly shows that, for the same ZrB2 and SiC starting powders and processing conditions, the density increases with increasing vol% SiC. This is reflected in the SEM micrographs presented in Fig. 4A-D. The grain sizes (of both ZrB2 and SiC) are similar between these four materials (Table 1), which is also evident from the SEM micrographs in Fig. 4A-D. There appears to be some coarsening of ZrB2 , where the starting-powder size is 2 \\u242em, and the measured grain sizes are between 3.4 and 3.9 \\u242em. Little coarsening is observed in SiC, which is to be expected considering the low hot-pressing temperature of 1650 C and the discontinuous nature of the SiC phase.     Fig. 4. SEM micrographs of materials: (A) Z-1, (B) Z/S-1, and (C) Z/S-2 showing the effect of increasing SiC content (same SiC starting-powder size of 1.7 \\u242em)  on the microstructures, and (D) Z/S-3, (E) Z/S-4, and (F) Z/S-5, showing the effect of decreasing SiC starting-powder size (same SiC content of 22.4 vol%) on the  microstructures. The gray phase is ZrB2 and the dark phase is SiC.  \\x0c', '220  S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  made from 40-nm SiC starting powder, showing the distribution of SiC grains. For material Z/S-6, where the 40-nm SiC powder was colloidally dispersed, the highest density of 99.9% was obtained (Table 1; Fig. 3B), and the SEM micrograph in Fig. 5B shows smaller SiC grain size and better dispersion. Although the SiC grain sizes in materials Z/S-5 and Z/S-6 appear to be 0.6 \\u242em in the SEM, TEM studies (Fig. 6A) reveal that the SiC exists as aggregates of many nanograins in those materials. The these SiC nanograins is estimated at 80 nm. average size of In contrast, the material Z/S-3 (made from 1.7-\\u242em SiC starting powder) contains single-crystal SiC grains, as seen in the TEM micrograph in Fig. 6B. In both cases the ZrB2 (hexagonal) and SiC (cubic) were confirmed by selected area diffraction patterns (SAEDP) in the TEM. Comparison of densities of materials Z/S-3 and Z/S-7 in Table 1 and Fig. 3A shows that, much higher densities are achieved in materials Z/S-3 (97.9%) made from ball-milled powders, in contrast to material Z/S-7 (90.9%) made with nonball-milled powders. The density improves to 96.4% in material Z/S-8 made with non-ball-milled powders but with ZrO2 addition. Fig. 7A is a bright-field TEM micrograph of material Z/S3, showing, in addition to ZrB2 and SiC, a Zr(O,B)2 phase. The amount of this phase is quite small, which is perhaps why it is not detected in the XRD pattern (Fig. 2). The qualitative elemental composition of that phase was confirmed in the TEM using EDS (Fig. 8B) and EELS (not shown here). The Cu and the Y in the EDS spectrum in Fig. 7B is from the TEM grid and the YTZ ballmilling media, respectively. It appears that the Zr(O,B)2 phase is most likely a ZrO2 -rich substitutional solid-solution of ZrO2 and ZrB2 , with a distorted tetragonal structure. This structure was revealed by SAEDPs collected from the Zr(O,B)2 region during careful tilting experiments in the TEM. Fig. 7C is an example of such a SAEDP using [110] zone axis, showing bro Fig. 5. SEM micrographs of materials: (A) Z/S-5 and (B) Z/S-6 showing the  effect of colloidal processing on the microstructure. The gray phase is ZrB2 and the dark phase is SiC.  A comparison of density data in Table 1 and Fig. 3B for materials Z/S-3, Z/S-4, and Z/S-5 shows that, for the same vol% SiC (22.4 vol%) and processing conditions, the density increases slightly with decreasing size of SiC starting powders. While the ZrB2 grain size in these three materials is similar (3.4-3.9 \\u242em), the SiC grain size is seen to decrease with decreasing SiC starting-powder size (Table 1 and Fig. 4D-F). Fig. 5A is a higher magnification SEM micrograph of the material Z/S-5,  Fig. 6. Bright-field TEM micrographs of materials: (A) Z/S-6 and (B) Z/S-3. In (A), the SiC aggregates of 0.6 \\u242em in size consist of 80-nm SiC grains, while in (B) SiC is not aggregated. The inset in (B) is SAEDP from a SiC grain using zone axis B = [1 1 0].  \\x0c', 'S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  221  Fig. 7.  (A) Bright-field TEM image of materials Z/S-3 showing, in addition to ZrB2 and SiC grains, a Zr(O,B)2 phase. (B) EDS spectrum from the Zr(O,B)2 phase. (C) SAEDP from the Zr(O,B)2 phase showing tetragonal distortion.     ken symmetry; the orthogonality between horizontal and vertical spot-rows is off by 2 . Zr(O,B)2 substitutional solid-solution phase has been observed before, but that phase was ZrB2 -rich, with a monoclinic structure [12]. TEM studies of the material Z/S-7 made from non-ball-milled powders did not show any evidence of the Zr(O,B)2 phase. Fig. 8 is a typical TEM micrograph from material Z/S-7 showing only the ZrB2 and SiC phases. Fig. 9A and B are high-resolution  TEM images of typical interphase interfaces between ZrB2 and SiC in materials Z/S-3 and Z/S-7, respectively. While the Z/S-3 material made from ball-milled powders has amorphous films at most of the interphase interfaces (Fig. 9A), the Z/S-7 material made from non-ball-milled powders appears to have less incidence of amorphous films at interphase interfaces (Fig. 9B). This indicates that the ball-milling process introduces conditions favorable for the formation of the Zr(O,B)2 phase and amorphous interphase interfaces.  3.2. Oxidation  Fig. 10A-D are cross-sectional SEM micrographs of oxidized, dense ZrB2 /22.4 vol% SiC materials (Z/S-3, Z/S-4, Z/S-5, and Z/S-6). The oxidation behaviors of all these materials appear to be similar, and consistent with what has been observed by others for ZrB2 /SiC ceramics [4,6], except for the thicknesses of the oxide layers. It has been shown that the outermost layer (top) is silica glass, designated as layer I. The compositional maps in Fig. 11A and B confirm the layer I to be rich in Si and O. The layer below that is a composite of ZrO2 /silica, which is somewhat depleted in Si, as seen in the compositional map in Fig. 11A. This depleted layer is designated as layer II. Below layer II is the unoxidized ZrB2 /SiC base material. The micrographs in Fig. 10A-D, together with the quantitative data reported in Table 2, show clearly that SiC grain reduction results in a dramatic reduction in the thicknesses of the oxide layers in these materials. The largest decrease in the thicknesses of layers I and II occurs by going from a SiC grain size of 1.8-0.8 \\u242em. With further decrease in the SiC grain size, the decrease in the layer thicknesses is marginal, with the colloidally processed material  Fig. 8. Bright-field TEM micrograph of material Z/S-7, where the Zr(O,B)2 phase could not be found.  \\x0c', '222  S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  Fig. 9. High-resolution TEM micrographs of typical interphase interfaces (arrows) in materials: (A) Z/S-3 and (B) Z/S-7.  Fig. 10. Cross-sectional SEM micrographs of oxidized materials (1500     C, 10 min, in air): (A) Z/S-3, (B) Z/S-4, (C) Z/S-5, and (D) Z/S-6. The oxidized surface is  at the top. All main images were taken at the same magnification. Insets are SEM micrographs of near-surface regions at higher magnification (same for all insets).  The silica layer I and the depleted sub-layer II are defined in (A).  Fig. 11. Elemental composition maps of cross-section of oxidized material Z/S-4 in SEM/EDS: (A) Si and (B) O. The oxidized surface is at  the top. Both images  were taken at the same magnification.  \\x0c', 'S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  223  4. Discussion        The results presented in Table 1 and Fig. 3A show clearly that the addition of SiC particles promotes densification of ZrB2 ceramics at a moderate hot-pressing temperature of 1650 C. The earlier published studies appear to have overlooked this fact, reason for which is not entirely clear. It is worth noting that in most of those studies dense, pure ZrB2 is used as a reference material, which requires high temperature hot-pressing (1900-2100 C). It is possible that, in an effort to maintain consistent processing conditions, the same high hot-pressing temperatures were used to densify ZrB2 with SiC additions. Also, while some studies report the use of WC-Co balls media [7-9], other do not report if ZrO2 ball-milling media was used. The explanation for the beneficial effect of SiC on densification may lie in the fact that most SiC powder particles are known to have an oxidized surface layer. That layer may promote formation of liquid phases during hot-pressing, assisting in densification at lower temperatures. Improvement in densification with increasing SiC content (Fig. 3A), and hence increasing oxide content, supports this hypothesis. Evidence of amorphous films at interphase grain boundaries (Fig. 9A) also lends support to this hypothesis. The observed increase in density of ZrB2 /22.4 vol% SiC materials with decreasing SiC starting-powder size (Fig. 3B) further supports this hypothesis: the higher surface oxide content in smaller SiC particles (high specific surface area) is expected to promote densification. The Z/S-7 material made from non-ball-milled powders has lower density (90.9%) compared to the material Z/S-3 made from ball-milled powders (97.9%). It may be argued that ballmilling results in better mixing and particle-size reduction, which in turn results in better densification. However, that is not a complete explanation because, as mentioned earlier, some ZrO2 from the YSZ ball-milling media gets incorporated indirectly in the ZrB2 /SiC powders during ball-milling. Direct addition of ZrO2 powder to the non-ball-milled ZrB2 /SiC powder mixture results in a density of 96.4% (material Z/S-8), which is in-between densities of materials Z/S-7 and Z/S3. This shows that ball-milling, together with the attendant ZrO2 addition, results in high densities in ZrB2 /SiC materials. Once again, the added oxide appears to promote densification during hot-pressing via introduction of liquid phases. It is not clear what role of the Zr(O,B)2 phase plays in the densification. There is clear evidence that decreasing the SiC grain size in ZrB2 /SiC materials results in an increase in the oxidationresistance. A qualitative explanation for this beneficial effect of SiC grain reduction lies in the consideration of the beneficial effect of any SiC additions to ZrB2 . It is generally accepted that the oxidation of pure ZrB2 proceeds with the formation of ZrO2 and boria at the exposed-surface. Boria is highly volatile, and it does not provide adequate protection against further oxidation of ZrB2 . In the case of ZrB2 /SiC materials, the oxidation of SiC grains provides a steady supply of silica glass. The silica glass combines with the boria formed on the neighboring ZrB2 grains, to result in a more refractory borosilicate glass. With further evaporative loss of boron, the borosilicate glass eventually  Fig. 12. SEM micrograph of oxidized top surface of material Z/S-4.  (Z/S-6) showing highest oxidation-resistance of all materials evaluated in this study. Fig. 12 is a SEM image of the top surface of material Z/S-4 showing smooth silica-rich glaze.  the oxidized  3.3. Mechanical properties  The mechanical properties results for the dense ZrB2 / 22.4 vol% SiC materials (Z/S-3, Z/S-4, Z/S-5, and Z/S-6) are reported in Table 2. Both the hardness and toughness are found to increase with decreasing SiC grain size and better dispersion. Fig. 13A and B shows microstructural interactions of indentation cracks in materials Z/S-3 and Z/S-6, respectively. The crack in material Z/S-3 with the coarser SiC grains (Fig. 13A) appears to be wavy and heavily bridged by the SiC grains, while the crack in material Z/S-6 with finer SiC grains (Fig. 13B) appears relatively straight with little or no bridging.  Fig. 13. SEM micrographs  showing interaction of  indentation cracks with  microstructure in materials: (A) Z/S-3 and (B) Z/S-6. The gray phase is ZrB2 and the dark phase is SiC. Arrows in (A) point to SiC-grain crack-bridging sites.  \\x0c', '224  S.S. Hwang et al. / Materials Science and Engineering A 464 (2007) 216-224  becomes silica-rich glass, which essentially glazes the surface and provides oxidation protection. Now consider a decrease in the size of the SiC grains from 1.8 to 0.8 \\u242em, while maintaining the SiC content of 22.4 vol%. Assuming spherical SiC grains and uniform distribution, this results in an increase in the ZrB2 /SiC interface-length per unit area of exposed-surface and a decrease in the spacing between SiC grains, both by about a factor of 2. This is expected to make SiC more effective in providing the silicarich glass, leading to the formation of the protective silica-rich ment, a further decrease in the SiC grain size to 80 nm is layer on the material early on during oxidation. Using this arguexpected to result in a dramatic improvements in the oxidationresistance in materials Z/S-5 and Z/S-6. However, only marginal improvements in the oxidation-resistance in materials Z/S-5 and Z/S-6 have been observed (Fig. 10 and Table 2). This individual 80-nm SiC is because, as seen in Fig. 6A, the form aggregates of 0.6 \\u242em size, grains in these materials which is the effective grain size in the context of oxidation. These results suggest that further improvements in the oxidation-resistance of ZrB2 /SiC materials could be realized by obtaining better dispersion of the SiC nanodispersoids in ZrB2 . The hardness of the dense ZrB2 /SiC materials reported in Table 2 increases with decreasing SiC grain size. This can be attributed to dispersion hardening of ZrB2 with the harder SiC phase that is finer and better dispersed. The higher toughness of the dense ZrB2 /SiC materials with coarser SiC grains can be attributed to crack-bridging toughening by SiC grains (Fig. 13A), which is more effective in coarser microstructures than finer microstructures [13].  5. Summary  • Addition of SiC promotes densification of ZrB2 at a moderate hot-pressing temperature of 1650 C; the highest density achieved in ZrB2 /SiC materials is 99.9%. A reduction in the SiC starting-powder size, and colloidal dispersion of the powders, both promote hot-pressing densification of ZrB2 /SiC materials.     • Ball-milling of the ZrB2 /SiC starting-powder mixtures using YSZ balls media results in the doping of the powder mixture with ZrO2 , which promotes hot-pressing densification. The resulting ZrB2 /SiC materials contain a Zr(O,B)2 phase, and amorphous films at interphase interfaces. • SiC grain-size reduction results in improved • The ZrB2 /SiC materials produced here possess modest hardresistance in ZrB2 /SiC materials. ness and toughness properties. • The results presented point to a new strategy for improving processing and oxidation-resistance of ZrB2 /SiC materials: dispersion and reduction of SiC grains.  oxidation Acknowledgements  The authors thank Profs. J. Li and R.A. Rapp for fruitful discussions. Funding for SSH was provided by a Korean Research Foundation post-doctoral fellowship (KRF-2005-D00218).  References  [1] M.M. Opeka,  I.G. Talmy,  J.A. Zaykoski,  J. Mater.  Sci.  39  (2004)  5887-5904.  [2] A. Bongiorno, C.J. Forst, R.K. Kalia, J. Li, J. Marschall, A. Nakano, M.M.  Opeka, I.G. Talmy, P. Vashishta, S. Yip, MRS Bull. 31 (2006) 20-26.  [3] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, J. Eur.  Ceram. Soc. 19 (1999) 2405-2414.  [4] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, J.  Eur. Ceram. Soc. 22 (2002) 2757-2767.  [5] F. Monteverde, S. Guicciardi, A. Bellosi, Mater. Sci. Eng. A346 (2003)  310-319.  [6] E. Opila, S. Levine, J. Lorincz, J. Mater. Sci. 39 (2004) 5969-5977.  [7] W.G. Fahrenholtz, G.E. Hilmas, A.L. Chamberlain, J.W. Zimmermann, J.  Mater. Sci. 39 (2004) 5951-5957.  [8] A.L. Chamberlain, W.G. Fahrenholz, G.E. Hilmas, D.T. Ellerby, J. Am.  Ceram. Soc. 87 (2004) 1170-1172.  [9] A.L. Chamberlain, W.G. Fahrenholz, G.E. Hilmas, D.T. Ellerby, Refract.  Appl. Trans. 1 (2005) 1-8.  [10] B.R. Lawn, Fracture of Brittle Solids, second ed., Cambridge University  Press, Cambridge, U.K., 1993.  [11] D.B. Marshall, T. Noma, A.G. Evans,  J. Am. Ceram. Soc. 65 (1982)  C175-C176.  [12] C.F. Feng, L. Froyen, Acta Mater. 47 (1999) 4571-4583.  [13] N.P. Padture, J. Am. Ceram. Soc. 77 (1994) 519-523.  \\x0c']"
},{
  "_id": 100,
  "PDF": "In situ Formation of Oxidation Resistant Refractory Coatings on SiC-Reinforced ZrB2 Ultra High Temperature Ceramics.pdf",
  "Text": "['In situ Formation of Oxidation Resistant Refractory Coatings on  SiC-Reinforced ZrB2 Ultra High Temperature Ceramics  ‡,†,* Eugenio Zapata-Solvas, Daniel Doni Jayaseelan,  ‡  §,* and William E. Lee Peter Brown,  ‡,**  ‡  Centre for Advanced Structural Ceramics (CASC), Department of Materials, Imperial College, London SW7 2AZ, U.K.  §  Dstl, Porton Down, Salisbury Wiltshire, SP4 0JQ, U.K.  In situ oxidation resistant and solid refractory coatings have  been  generated  on  20 vol% SiC-reinforced ZrB2 temperature ceramics containing 10 wt% rare earth (RE) addi ultra  high  tives such as LaB6, La2O3, and Gd2O3 fabricated by spark plasma sintering. Oxidation for 1 h at 1600°C in static air led (up to 250 lm thick) of ZrO2 and RE-zirconates on the composite systems underneath which (50-100 lm)  to formation of a dense  layer  were  intermediate  layers  containing  heteroge neous crystalline oxides such as La2Zr2O7 and amorphous silicate phases. The benefits of predominating solid oxidation  products  over  having  substantial  volumes  of  liquid  present  in  aerospace leading edge applications are discussed.  I.  Introduction  U LTRA high temperature ceramics (UHTCs) are candidate materials for a variety of aerospace applications owing to their unique combination of properties, including high melting temperature (>3000°C), high strength, and high elastic modulus.1-5 However, issue in the development of UHTCs  oxidation  resistance  is  a major  for aero-propulsion and  hypersonic  flight  applications. The (Me = Zr monolithic MeB2 and Hf) ceramics and B2O3. Below 1200°C, formation of MeO2 liquid B2O3 fills the pores between oxide grains, providing oxidation protection. However, above 1200°C, MeB2 oxidizes result of volatilization of the protective B2O3 in the oxide scale6-9 which leaves behind a porous nonprotective columnargrained ZrO2 or HfO2 layer. Much effort has been aimed at improving the oxidation resistance of ZrB2and HfB2-based composites6,10-25 focusing on the addition of 20-30 vol% SiC. Compared with liquid  oxidation  resistance  of  depends  on  rapidly as a  B2O3, borosilicate (BS) glass has higher melting temperature, higher viscosity, and lower oxygen diffusivity, thus providing more effective oxidation protection.26,27 SiC-containing ZrB2 ceramics have relatively good oxidation resistance to 1500°C. Above 1200°C, tion resistance by formation of a prophylactic BS glass coatHalloran32  the addition of SiC provides improved oxida ing.10-14,28-32  Karlsdottir  and  used  high  temperature optical microscopy to reveal that convection cur rents were set up in the BS liquid leading to “volcanoes” and  burst  bubble  structures  on  the  surface of the (>1800°C)  oxide  scale.  However, at very high temperatures  the BS/SiO2 to continuous  liquid is  lost  exposing  the underlying boride  active oxidation. As a result there is much interest in develop ing  composite  systems with  improved  oxidation  resistance  above  this  temperature. To  further  improve  the  oxidation  resistance  of  ZrB2/SiC, CrB2, TiB2, TaB2, NbB2 and VB2, have been added to the MeB2-SiC composite system.33 Improved oxidation resistance was related to the presence of transition metal oxides in the  transition metal  borides,  such  as  BS inducing  its phase  separation (immiscibility)  leading  to  increased liquidus temperatures and viscosities. These charac teristic  features  of  immiscible  glasses,  are  beneficial  for  decreasing oxygen diffusivity and suppressing boria evaporation from the glass.34 Others26,33-36,37 have studied the effect of Si-containing additives, such as Si3N4 and silicides, on the oxidation behavior of ZrB2 and HfB2 ceramics. Talmy et al.26 discussed the role of different silicides on the oxidation behavior of ZrB2. However, Opila et al.38 reported that addition of 20 vol% TaSi2 to ZrB2/20 vol% SiC improved its oxidation resistance at 1627°C in air. The improved behavior is attributed to the presence of Ta, not to the increased Si  content  arising from the additional Si-containing component. In gen eral,  the oxidation behavior of non-oxide  ceramics depends  largely on the chemical composition and properties of the oxi dation  products  and  on  the  combination  of  physical  and  chemical processes taking place on the oxygen atmosphere.26 Modification of  surface  exposed to  the chemical composi tion of  the oxide layer,  leading to decreased inward diffusion  of oxygen,  is an effective way of controlling oxidation resis tance of non-oxide ceramics. This modification can be accom plished et al.39  by  changing  the  bulk  ceramic  composition. Zhang  added WC to ZrB2 (LPS) ultimately leading to formation of a two-layer dense,  inducing  liquid phase  sintering  protective,  oxide  scale;  the  outer  layer  comprising  ZrO2,  B2O3, and WO3 and the inner layer ZrO2 and WO3. There have been many previous attempts to improve  the  oxidation  resistance  of  ZrB2/HfB2 the viscosity of the liquid silica layer,  ceramics  by  for  e.g.  increasing  increasing  the  immiscibility of different  liquid phases, or by develop ment of dense ZrO2 layers via LPS thus penetration and diffusivity.40 However,  reducing  oxygen  in hypersonic  space  and military 2000°C in temperatures and in such harsh environments,  applications,  the  temperature  often  exceeds  combination with  severe  airflow. At  such  high  the BS melt  will be blown off rapidly exposing the underlying layer. This  may lead to further oxidation and UHTC material recession.  This work aims  to develop solid refractory oxide protective  layers by  adding  rare  earth (RE) borides or oxides  to the  baseline (ZrB2/20 vol% SiC) UHTCs. Refractory oxide layers will be formed by in situ reaction40 during the oxidation  process.  II.  Experimental Details  (>99%, ~45 lm, ρ = 6.85 g/cm3; ZrB2 powder SigmaAldrich, Gillingham, U.K.) and SiC powder (a-SiC, 99%, d50 ~0.7 lm, ρ = 3.217 g/cm3; Good Fellow Chemicals, Hunting d50  J. Halloran—contributing editor  Manuscript No. 29589. Received April 13, 2011; approved November 22, 2011.  *Member, The American Ceramic Society  **Fellow, The American Ceramic Society  Crown copyright 2011. Published with the permission of  the Defence Science and  Technology Laboratory on behalf of  the Controller of HMSO.  †  Author to whom correspondence should be addressed. e-mail: d.j.daniel@imperial.  ac.uk  1247  J. Am. Ceram. Soc., 95 [4] 1247-1254 (2012)  DOI: 10.1111/j.1551-2916.2011.05032.x  Published 2012 by Wiley Periodicals, Inc.  Journal  \\x0c', 'don, U.K.) were used as starting materials to form a baseline  ZrB2/20 vol% SiC (hereafter ZrB2 (space group, P6/mmm) P63mc) have a hexagonal structure with lattice parameters of a = 3.17 A˚ , c = 3.53 A˚ and a = 3.08 A˚ , c = 15.12 A˚ , respec termed  ZS20) UHTC.  Both  and  6H-SiC (space  group,  tively. To  improve  the  oxidation  resistance (>99%, (>99%,  of ZS20, RE ~2-3 lm, ~2 lm,  additions in the LaB6 ρ = 4.72 g/cm3; Sigma-Aldrich), La2O3 ρ = 6.51 g/cm3, Fluka Chemicals supplied through SigmaAldrich, Steinheim, Germany) and Gd2O3 (>99%, d50 ~2 lm, ρ = 7.07 g/cm3, Fluka Chemicals supplied through Sigma form  of  d50  d50  Aldrich) were added to the starting materials. The as-received  ZrB2 powder was mill” shatter box using a steel  further dry-milled for 30 min in a “swing container  coated with Teflon  to reduce the average particle size. The average particle size  (measured  using  a Malvern  ®  laser-diffraction unit) of the 5.11 ± 0.5 lm. Longer in significant reduction  powder obtained after milling was milling times (>30 min) did not result  in the particle  size. Appropriate amounts of ZrB2, SiC, and RE additive were wet ball milled in a plastic container using  ZrO2 rotary  balls  in  ethanol  for  nearly  12 h  and  dried  using  a  evaporator. The dried powder  compositions, namely  ZS20/10 wt% La2O3 10 wt% LaB6 (hereafter 10 wt% Gd2O3 fied in a spark plasma sintering (SPS)  (hereafter  referred to as ZSLO), ZS20/  referred  to  as ZSLB)  and ZS20/  (hereafter  referred to as ZSGO) were densi furnace (FCT Systeme,  GmbH, Rauenstein, Germany). A 20-mm diameter graphite  foil-lined die was used. The  graphite die was  covered with  graphite  felt  to reduce heat  loss and the  temperature moni tored by an optical pyrometer which was  sighted from the  top of the graphite punch. Samples were sintered under 5 Pa vacuum between 1750°C and 1850°C for less than 10 min. A 100°C/min was maintained to the heating rate of sintering ~70 MPa was  temperature  and an applied load of  applied  during sintering. Bulk density measurements were carried out  using the Archimedes method in water. Relative density was  calculated by dividing measured bulk density by theoretical  density (TD) calculated by the rule of mixtures.  The RE-doped  samples,  20 mm in  diameter  and  5 mm  thick were cut  in half and placed on an alumina boat  so as  to have minimum point of contact between the sample and boat and oxidized at 1600°C for 1 h in static air using a lab oratory  open  hearth  furnace. Oxidized  samples were  cut,  mounted in epoxy resin molds, and ground and polished using various grades of media down to 1 lm. For compari(~99% of TD) monolithic ZrB2 and ZS20 were fabricated using SPS at 1900°C and 1850°C, identical heating rate, pressure and  son, dense  also  respec tively, with application of  holding  time,  and oxidation tests were performed on them  under the same conditions.  Phase analysis of  sintered and oxidized samples was  car ried out using X-ray diffraction (XRD) on a Philips PW7100  Diffractometer  (Eindhoven,  the Netherlands)  using  CuKa  radiation.  International  center  for diffraction  data  (ICDD)  cards were used to identify phases ZrB2 SiC (00-029-1131), Gd2O3 (01-074-1987), La2O3 1345), LaB6 (34-0427), La2Zr2O7 (17-0450), Gd2Zr2O7 0799), t-ZrO2 (00-048-0224), and m-ZrO2 (00-037-1484).Thermal analysis of sintered samples was carried out using a ther (01-075-1050), 6H (01-083 (16 mal  analyzer  (Netzsch STA 449F1, Wittelsbacherstrasse, from room temperature to 1600°C at a rate  Selb, Germany) 10°C/min samples were  of  under  flowing  air  (50 mL/min). At  least  two  tested in each composition using  the  thermal  analyzer under identical conditions and no significant change  was observed between them. Plan view and cross  sections of  the  polished  sintered  and  oxidized  samples were  observed  using  a  scanning  electron microscope  (SEM)  fitted with  a  field-emission gun (FEG, model LEO15). Secondary electron  (SEI) and back-scattered electron images (BEI) were taken at an operating voltage of 20 keV, and a beam current 105 lA  and at a working distance between 10 and 15 mm. Chemical  analyses were carried out using an energy dispersive spectros copy  (EDS)  unit  (Oxford  Instruments, Oxfordshire, U.K.)  attached to the FEG-SEM. For all samples, tance was ~10 mm. imaged and the oxidized  the working dis In addition,  the oxidized samples were  layers were  compositionally  ana lyzed in an FEI FIB-SIMS 200T focused ion beam work sta tion  (FIB) with  secondary  ion mass  spectroscopy  facility  attached (SIMS). Gallium ions were used to bombard the  sample  surface. The  samples were attached to an aluminum  sample holder with silver  tape  to avoid charge build-up on  the  sample  surface.  Specimens  for  transmission  electron  microscopy  (TEM)  observations were  prepared  from SPS  materials  using  conventional mechanical  polishing  and  ion  thinning.  Ion thinning was performed using a Gatan Model  691 precision ion polishing  system (PIPS), Gatan, Oxford shire, U.K. TEM sections on the oxide layers were prepared  using a FIB work station operating with a gallium beam at  30 keV. Bright-field (BF)  images and selected area electron  diffraction  (SAED)  patterns were  acquired  using  a  JEOL  JEM-2000EX  (Oxfordshire,  UK)  transmission  electron  microscope operating at 200 kV with an Oxford Instruments  EDS microanalysis  system and SAED patterns were  solved  using the ratio method.  III.  Results  All samples, namely ZS20 sintered at 1800°C, ZSLO sintered at 1850°C, ZSLB sintered at 1750°C, and ZSGO sintered at 1800°C, Phase analysis  attained  above  99% of TD in  less  than  10 min.  (Fig. 1)  confirmed that all  starting materials  remained 1600°C,  after  SPS. However,  after  oxidation  for  1 h  at  phase  transformation  and  in  situ  reactions  have  occurred  on  the  exposed  surfaces  of  these  samples. XRD  detected  predominantly m-ZrO2 samples after oxidation at  in monolithic ZrB2 1600°C. Figure 2 shows  and  ZS20  XRD of  oxidized  surfaces  of  ZSLO,  ZSLB,  and  ZSGO  revealing that  the  sample  surfaces  comprised predominantly  m-ZrO2. A trace of t-ZrO2 was also observed. In addition, situ formation of RE zirconate (RE2Zr2O7) pyrochlore was observed during oxidation. Figure 3 shows ground and pol in  ished  surface microstructures  of  ZSLO [Fig. 3(a)],  ZSLB  [Fig. 3(b)],  and  ZSGO [Fig. 3(c)]  after  SPS.  The  surface  microstructure of ZS20 consisted of a homogeneous distribution of SiC in ~5 lm diameter ZrB2 grains described in a (4.0-7.0 lm) medprevious article.41 Figure 3 reveals large ium gray ZrB2 grains, finer (~1.0 lm) angular dark gray SiC grains, and light gray (>2.0 lm) RE additive phases. In all  three  compositions, La2O3, LaB6, and Gd2O3 grains always appear agglomerated together with SiC forming an intercon nected  network  and  are  distributed  homogeneously  in  the  Fig. 1.  XRD of  spark  plasma-sintered  ZS20  with  various RE  additives; no phase transformation observed.  1248  Journal of  the American Ceramic Society—Jayaseelan et al.  Vol. 95, No. 4  \\x0c', 'April 2012  Formation of Oxidation Resistant Refractory Coatings  1249  Fig. 2.  XRD of ZS20 containing La2O3 (ZSLO), LaB6 (ZSLB) and (ZSGO) oxidized for 1 h at 1600°C showing product oxides  Gd2O3 and in situ formed zirconate phases.  (a)  (b)  (c)  Fig. 4.  Bright-field  TEM image  of  ZSGO showing  a  range  of  phases. EDS from the labeled grains is shown.  ZrB2 matrix. Figure 4 shows a bright-field (BF) sintered ZSGO which is representative of all three composi image of as tions. Different  grains  are  labeled  and  the  corresponding  shown.  EDS are  In addition to the ZrB2, SiC, and Gd2O3 starting materials, other phases present are ZrO2 (grain A) possibly arising from grinding media and grain E in contact  with SiC is a Gd-based silicate phase whose morphology sug gests that it is glassy.  Figure 5(a) shows SEI of a plan view of a ZrB2 sample after oxidation for 1 h at 1600°C revealing the porous nature  of the surface. EDS detects only Zr and O. In cross [Fig. 5(b)], a 40-lm thick oxidized porous ZrO2 revealed on the unaffected ZrB2. Figure 6 shows plan view and cross-section microstructures of baseline-sintered ZS20 after oxidation for 1 h at 1600°C.  case of ZS20,  top layer  In the  section  is  the  surface features and nature of the oxidized layers are differreveals fine (1-2 lm) bright contrast  ent. Figure 6(a)  spheri cal particles in a dark gray matrix. These are secondary ZrO2 particles which most likely precipitated from liquid silicate.20,28 The cross-section microstructure of ZS20 [Fig. 6(b)] shows two layers, an outer 5 lm thick silicate layer containing ZrO2 particles and an intermediate ~30 lm thick porous ZrO2 layer between the outer silicate layer and the underlying unaffected ZrB2-SiC region. Figure 7 shows the surface and cross-section microstructures of ZSLO oxidized 1 h at 1600°C. The oxidized surface (1 lm) and grain bound [Fig. 7(a)] contains  spherical pores  ary  cracks. The  smooth and rounded nature of  the micro structural  features  suggests  liquid  formation  during  oxidation. EDS [Fig. 7(b)]  from the oxidized surface  reveals  the presence of Zr, La, Si, and O in the liquid. The cross sec tion [Fig. 7(c)] ~250 lm thick  shows  two distinct oxidized layers. The outer  layer  is  predominantly  a  continuous  bright  phase with isolated dark regions present. EDS of  the circled  region 1 of bright phase  in the outer  layer  reveals La, Zr,  and O suggesting the presence of La2Zr2O7 as confirmed by XRD (Fig. 2). Figure 7(d) shows BSI of the dark area of the  circled region 2 in Fig. 7(c) consisting of two phases, possibly  arising from phase  separation of  isolated droplets of  liquid.  EDS of circled region 2 reveals  the presence of La, Si, and O.  Fig. 3.  Microstructures of as-fabricated ZS20 with RE additives (a)  BEI of ZSLO, (b) BEI of ZSLB, and (c) BEI of ZSGO.  \\x0c', 'The  rough intermediate layer above the unaffected bulk is ~50 lm thick, EDS reveals that Zr and O are the main elements in it.  material  Figure 8 shows a region of the electron image of the oxiin ZSLO obtained at a tilt angle of 45°  dized outer  layer  in  the FIB workstation. The  electron image  [Fig. 8(a)]  shows  bright and dark phases. SIMS analysis [Fig. 8(b)] was carried  out on some grains causing the craters in Fig. 8(a). SIMS clearly reveals the presence of La, Zr, and Zr-O in the bright  grains  suggesting that  they are  lanthanum zirconate grains.  Figure 8(c) shows a BF image of a TEM section of a sample  ion-milled from the interface of  the outer  layer and interme diate layer of  the oxidized region consisting of at  least  three  different phases, namely La2Zr2O7, ZrO2, and silicate glass. Figure 8(d) shows the SAED pattern of circled region d in  Fig. 8(c)  tilted to the  [001]  zone  axis. The pattern can be  indexed as pyrochlore cubic structure with a lattice parameter = 10.768 A˚  corresponding  to  La2Zr2O7. in Fig. 8(c). La, Zr, and  Figure 8(e)  shows  the EDS taken on region ‘d’  O are  its main  constituents  suggesting  the  formation  of  La2Zr2O7. Figures 9(a) and (b)  show the microstructures of  the oxi dized surface and Figs. 9(c) and (d) show details from a cross  section of ZSLB. The oxidized surface is heterogeneous and  covered with  a  variety  of  phase morphologies.  In  some  regions  [Fig. 9(b)],  it  appears  to  be  compact with  grain  boundary phases presumably derived from liquid. Large pits ~600 lm in  diameter  are  formed  on  oxidation  [Fig. 9(a)].  Two layers can be seen at low magnification in Fig. 9(c). The outer layer is ~120 lm thick and dense while the intermediate ~50 lm thick. The  layer  is  continuous  phase  in  the  outer  layer contains mainly Zr, La, and O (EDS 1) while the iso lated dark regions  (EDS 2) are mostly silica suggesting that  they derive from liquid since silica would melt at  this temper ature. A flowerlike pattern is occasionally observed in the  (a)  (b)  Fig. 5.  ZrB2 oxidized for 1 h at 1600°C. porous ZrO2 microstructure and (b) BEI of two regions; porous ZrO2 top layer, I, and un-reacted ZrB2 Representative EDS from regions 1 in (a) and (b) is shown.  (a) SEI of surface showing  cross  section showing  layer.  (a)  (b)  Fig. 6.  SEIs of ZS20 oxidized for 1 h at 1600°C.  (a) Plan view of  surface showing bright contrast ZrO2 particles (I) and (b) Cross section showing three regions; SiO2 + ZrO2 intermediate porous ZrO2 reacted ZrB2/SiC region. EDS from regions shown.  (II)  in silicate matrix  top  protective  layer,  layer and bottom un I  and  II  in  (a)  are  1250  Journal of  the American Ceramic Society—Jayaseelan et al.  Vol. 95, No. 4  \\x0c', 'microstructure  of  the  cross  section  of  the  oxidized  layers.  Figure 9(d)  shows a BEI of  such a region in the  top outer  layer  in cross  section at higher magnification revealing sev(10 lm)  eral phases around the  light ZrO2 grains petal-like silica (EDS  including  pockets  of  dark  2)  decorating  the  ZrO2, phase  and  between  and  around  them a  continuous  bright  (EDS 3)  containing nanoscale dark particles. EDS 3  reveals that the bright phase is predominantly La and Si with  high O and  low Zr  presumably  from neighboring  grains  which along with its morphology suggests  it  is a lanthanum  silicate glass whereas the dark particles are silica glass. Clearly, La ions have concentrated at 1600°C to form lantha num silicate glass within the predominantly silica glass,  fur ther  evidence of phase  separation of  two immiscible glasses  in this study. EDS 4 is from the second layer and reveals the  presence of Zr, Si, La, and O. The layer below the top layer  is porous with a rough appearance.  Figure 10 shows  the top outer exposed surface and cross1600°C.  section microstructures  of ZSGO oxidized  1 h  at  SEI of  the outer  exposed surface  reveals  it  is  compact with  few shrinkage  cracks  forming  on  cooling  [Fig. 10(a)]. Fig ures 10(b) and (c) are SEI images of  the cross-section micro100 lm thick  structure  showing  three  phases  in  the  outer  oxidized layer with black, dark,  gray,  and bright  contrast.  EDS reveals that  the light gray contrast grains indicated by 1  are ZrO2, O, and Si and the black contrast grains indicated by 3 are predominantly silica. The intermediate layer is ~50 lm thick. Figure 11 shows thermogravimetric analysis of ZrB2, ZS20, and RE-ZS20 carried out under compressed air flow from room temperature to 1600°C. There was 5.14 ± 0.2% increase in mass for ZrB2 and 2.91 ± 0.1% for ZS20. LaB6 and La2O3 additions had 3.41 ± 0.2% and 4.12% ± 0.2 mass respectively, whereas Gd2O3 had only 1.8 ± 0.2% increases, mass increase.  the bright phase  indicated by 2 contains Gd, Zr,  IV.  Discussion  Ground and polished SEI surface microstructures of sintered  samples of ZS20, ZSLO, ZSLB, and ZSGO were similar.  In  (a)  (b)  (c)  (d)  Fig. 7.  SEIs of ZSLO oxidized for 1 h at 1600°C.  (a) area exposed  surface showing many cracks and pores.  (b) corresponding EDS, (c) top protective La-Zr-O  cross  section showing two oxidized layers;  and intermediate ZrO2 layers on the un-affected ZrB2/SiC and (d) BEI image of the dark circled in region 2 in (c) showing its bi-phasic  nature. EDS from regions 1 and 2 in (c) are shown.  (a)  (c)  (b)  (d)  (e)  Fig. 8.  ZSLO oxidized 1 h at  1600°C (a) FIB image of oxidized  layer  showing different phases,  (b) SIMS of bright grains  in (a),  (c)  BF-TEM image of a region in oxidized layer,  (d) SAED of grain d  (La2Zr2O7) on grain “d” in (c) showing the presence of Zr, La, and O.  in (c)  taken along the [001] zone axis and (e) EDS taken  (a)  (b)  (c)  (d)  Fig. 9.  ZSLB  oxidized  for  1 h  at  1600°C.  (a)  SEI  of  exposed  surface  showing  large  pit  and  cracks,  (b)  SEI  showing  smooth  surface;  (c) SEI of cross  section showing different  layers,  (d) BEI of  a region in top protective layer showing a flowerlike pattern, contain  three different phases. EDS were taken on regions marked (1-4).  April 2012  Formation of Oxidation Resistant Refractory Coatings  1251  \\x0c', '(0.7 lm) the RE-added samples the small SiC particles are always in close proximity with the small RE2O3 particles (2- 3 lm) as would be expected when mixing two small size range particles with one larger (ZrB2). Previous studies42 suggested that during LPS of SiC, La2O3 reacts with the surface SiO2 on SiC to form a lanthanum silicate phase. It seems likely therefore that at the high sintering temperature, ~1800°  C, used in the present  study at  least  some of  the RE oxide  reacts with the  surface SiO2 on SiC to form a liquid grain silicate phase. Figure 4 suggests that in ZSGO,  boundary  reaction has occurred with the Gd phase to form a Gd-based the presence of <1 vol% of ZrO2  silicate glass. Furthermore,  is  likely to arise  from the ZrO2 grinding media used during from the surface oxide impurities inherent to the  milling or  starting ZrB2. RE phases (La2O3, LaB6, and Gd2O3) were added to baseline UHTCs with the intention that they react  with ZrO2 to form RE-zirconates during oxidation. Although the retention of the starting materials without any reaction  occurring between them is  important,  this does not  rule out  participation of any reacted phases  such as RE silicate and  zircon in later oxidation reactions. expected during oxidation are29:  The main  reactions  ZrB2ðsÞ þ 5 2  O2ðgÞ ! ZrO2ðsÞ þ B2O3ðlÞ  (1)  B2O3ðlÞ ! B2O3ðgÞ  (2)  SiCðsÞ þ 3 2  O2 ! SiO2ðlÞ þ 3O2ðgÞ  (3)  When adding LaB6 expected to occur during oxidation.  to ZS20,  the  following  reactions  are  LaB6ðsÞ þ 11 2  O2ðsÞ ! 1 2  La2O3ðsÞ þ 3B2O3ðlÞ  (4)  ZrO2ðsÞ þ La2O3ðsÞ ! La2Zr2O7ðsÞ  (5)  When  adding RE2O3 occurs during oxidation.  to  ZS20,  the  following  reaction  ZrO2ðsÞ þ RE2O3ðsÞ ! RE2Zr2O7ðsÞ  (6)  SiO2ðsÞ þ RE2O3ðsÞ ! RE silicate glass  (7)  (where RE = La and Gd in this the pyrochlore structure and melting temperatures \\x15 2300°C study) RE-zirconates have (La2Zr2O7—2295 ± 10°C and Gd2Zr2O7—2450 ± 10°C). Furthermore, formation of RE2Zr2O7 is expansive, i.e. reaction (5) occurs during oxidation,  if  1 unit cell volume of ZrO2 þ 1 unit cell volume of La2O3 ! 1unit cell volume of La2Zr2O7 3 ðfor m-ZrO2 Þ þ 82:9 ˚A 3 ! 1262:37 ˚A  i.e.,140:62 ˚A  3  indicating ~ a fivefold increase in unit volume of  the reaction  product. Owing to this volume expansion, RE-zirconate for mation is  likely  to contribute  to filling pores  generated by  evaporation of volatile species such as B2O3. During oxidation, ZrB2 oxidizes to form ZrO2 and B2O3, where B2O3 later volatilizes leaving behind porous ZrO2 [Fig. 5(b)]. Similarly, SiC oxidizes to form SiO2 [Fig. 6(b)]. These oxide phases, ZrO2 and SiO2, are available with RE2O3 to form RE2Si2O7 or RE2Zr2O7. However, XRD shows only formation of RE2Zr2O7. It seems likely that RE silicate formed as a liquid at high temperature and  39  to react  cooled to a  glassy phase  [Figs. 4(e),  7(d),  9(d),  and 10(c)],  and hence was not detected by XRD. The predominance of  m-ZrO2 implies almost complete transformation from t-ZrO2 has occurred, which again involves nearly ~3 vol% expansion.  Comparison  of  the  cross  sections  of monolithic  ZrB2 ZS20  [Fig. 5(b)],  baseline  UHTC  ZrB2/20 vol% with RE additions, [Fig. 7(c) SiC,  [Fig. 6(b)],  and  ZS20  for  ZSLO] [Fig. 9(c)for ZSLB] and [Fig. 10(b)for ZSGO] after oxidation for 1 h at 1600°C reveals that ZS20 composi tions with added RE are likely to have improved oxidation.  However, oxidation in these systems is a complicated process  Fig. 11.  Thermogravimetric  analysis  of  (a) monolithic ZrB2, (d) ZSLO, and (e) ZSGO from room temperature  (b)  ZS20, (c) ZSLB, to 1600°C.  (a)  (b)  (c)  Fig. 10.  ZSGO oxidized  for  1 h  at  1600°C.  (a)  SEI  of  exposed  surface  showing  gray  and  dark  phases  (b) BEI  of  cross  section  showing different  layers with irregular  top layer  thickness,  and (c)  BEI  of  region  2  in  top  protective  layer  showing  three  different  phases. EDS were taken from regions marked (1-3).  1252  Journal of  the American Ceramic Society—Jayaseelan et al.  Vol. 95, No. 4  \\x0c', 'involving  combinations  of  competing  reactions  leading  to  both mass  gain  and  loss.  In  general, mass  gain  alone  (Fig. 11)  is  not  a  reliable measure  of  oxidation  resistance  since the overall mass change is a combination of mass gain  due  to formation of  condensed species and loss due  to for mation of gaseous and volatile  species. Nonetheless,  in this  system gain in mass  is  likely indicative of  significant oxida tion due to the predominance of ZrB2 oxidation to ZrO2. the case of monolithic ZrB2, a weight gain of over 5% occurred (Fig. 11) with thinner depth (40 lm) of oxidation [Fig. 5(b)] and no layer developed which is likely to provide  In  oxidation protection since the one  that  is present  is porous.  During oxidation of ZrB2 into ZrO2 and B2O3, liquid B2O3 is completely volatilized at these temperatures leaving only porous ZrO2 grains.29 According to  reaction  (1),  1  molar  mass  of  ZrB2 (123.22 g/  (112.84 g/mol) gives  rise to 1 molar mass of ZrO2 mol) and 1 molar mass of B2O3 (69.62 g/mol).  1 molar mass of ZrB2 ¼ 112:84 g/mol  1 molar mass of ZrO2 ¼ 123:22 g/mol  1 molar mass of B2O3 ¼ 69:62 g/mol  Now there arise two possible situations,  (1)  If ZrB2 no escape of B2O3, of 141%.  fully oxidizes  to ZrO2 and B2O3 and there is there should be a weight increase  (2)  On  the  contrary,  if B2O3 there should still be a weight increase of 8%.  escapes  during  oxidation,  Figure 11 reveals only 5.14% increase  in mass  suggesting  that incomplete ZrB2 oxidation has 1600°C. However, the porous nature of oxygen transport through the pore channels and further deg taken place after 1 h at  this  layer  facilitates  radation  of ZrB2 would be expected.  under more  severe  oxidation  conditions  Baseline UHTC (ZS20)  had  a  thin  (<5 lm) outer silica layer [Fig. 6(b)], which hinders inward diffusion and the mass gain was ~2.91% (Fig. 11). (1-3),  protective  oxygen  In ZS20,  according to reactions  there should be a mass gain of  140% for  the  complete  conversion of  the original  specimen  to condensed oxide phases and in which the volatile  species (1-3),  do  not  escape. However,  in ZS20,  besides  reactions  interaction between reaction products like silica, boria and zirconia giving rise to zircon, BS, and Zr-BS also need to be  considered. No attempt  is made herein to quantify the vol umes  of  these  phases  to  link with  the mass  and  volume  changes due  to the difficulty of quantifying  the  level of B  present using  the  chemical  analysis  techniques  available  to  us.  In  RE-added  ZS20,  the  reactions during oxidation reports26,27,33 detected for become more  complex. Previous  mation of  liquid BS glass, which is  less viscous  than liquid  SiO2. Phase separation of this glass has been reported to help delay oxygen transport. In a borosilicate or silicate glass con taining transition metal  cations,  the  tendency toward liquid  immiscibility is known to increase with increasing cation field  strength of  the  transition metal. This phase  separation has  been argued to result  in increased viscosity, which has been  correlated  to  reduced  oxygen  diffusion  rates. Nonetheless,  this protective silica is liquid, and hence could not be used at 1800°C in  temperatures  above  an  aerospace  environment  with air flow as  the liquid would be blown off.  In addition,  the  second ZrO2 transport to the underlying composite. The RE-added ZS20, however, had thick (up to 250 lm)  layer  is  highly  porous  enabling  oxygen  outer layers [Figs. 7(c), 9(c) and 10(b)] and weight increases up to 4.12% (Fig. 11) during oxidation for 1 h at 1600°C. A  schematic  diagram comparing  the  cross  sections  of mono lithic ZrB2, ZS20, and ZS20 with RE-additives Fig. 12. Among the RE-additives, ZSGO had  is  shown in  only  1.81%  mass increase and is protected by a thick (average thickness ~150 lm) oxide layer [Fig. 10(b)] when compared with the thicker [Fig. 7(c), up to ~250 lm] but  less protective (4.12%  weight increase, Fig. 11) layers on La2O3 LaB6 [(Fig. 9(c), ~125 lm), Fig. 11, 3.41% weight increase)]. The  and  outer  layers are dense and consist predominantly of crystal line  refractory oxides  such as ZrO2 and RE2Zr2O7 and localized silicate phase. Althoughthe mass increase  [Fig. 8]  in  RE-added ZS20 is higher  than that  in ZS20,  the  formation  of dense solid layers  is  likely to be more advantageous  than  the sole formation of liquid silica. This is because at atures above 1600°C and high threshold velocity,  temper liquid silica  will be blown off by  the  air flow of  the  atmosphere with  which it  is  in contact paving the way for  further attack of  the surface. Chemical composition and phase analyses carried  out using EDS and XRD (Fig. 2) of  the outer  layers  reveal  that  in  general  they  contain mainly  ZrO2, RE2O3 intermediate layers revealed  and  RE2Zr2O7. EDS analysis of RE, Zr, O, and Si. Hence it  the  is  likely that Re2Zr2O7 and silitop surface during oxidation.  cate phases  formed below the  A possible  reaction sequence during oxidation is: ZrB2 and SiC oxidize to ZrO2, B2O3, and SiO2 phases and CO2 gas is emitted. This leads to formation of BS liquid. Later ZrO2 glass18-20 dissolve in this borosilicate temperature, B2O3 volatilizes leaving behind ZrO2 and RE2O3 grains in a silicate melt and a competition exists between ZrO2 and RE2O3 to react with SiO2. This leads to the formation of both RE2Zr2O7 and/or silicate phase(s) along with unreacted ZrO2 and BS glass. The distribution of Zr in both the unreacted  and RE2O3 increasing the  grains  liquid  viscosity. At  high  material and the oxide scale is continuous and homogeneous  indicating  that  the  oxide  scale was  coherent  and  compact  even  though many  voids were  present  owing  to  shrinkage  cooling. The overall  effect of  the RE addition is  to signifi cantly alter  the  chemical  composition and crystalline nature  of phases forming on the top oxidized layer.  V.  Conclusions  In-situ oxidation resistant and refractory coatings have been  generated on spark plasma sintered 20 vol% SiC-reinforced  ZrB2 10 wt% of rare earth (RE) additives such as LaB6, La2O3, and Gd2O3. Oxidation for 1 h at 1600°C in static air led to formation of a dense surface layer (up to 250 lm thick) of  (ZS20)  ultra  high  temperature  ceramics  containing  ZrO2 and RE-zirconates. With melting 1600°C it solid throughout  points well  above  is expected that both phases once formed remained  the oxidation process. Conversely,  the oxi dized surface of ZS20 without RE-additives comprised a por Fig. 12.  Schematic  comparison  of  the  cross  sections  of  (a)  monolithic ZrB2, (b) ZS20, and (c) ZS20-10% Re2O3/ReB6 oxidized for 1 h at 1600°C.  April 2012  Formation of Oxidation Resistant Refractory Coatings  1253  \\x0c', 'ous ZrO2 silica which was  layer  (up to 10 lm thick) covered by amorphous liquid at 1600°C. Owing to the low oxygen  permeability of  liquid silica its presence suppresses excessive 1600°C. However,  oxidation  of ZS20  in  static  air  at  in  a  hypersonic air  stream this protective advantage would likely  be quickly lost owing to liquid silica removal by viscous flow.  This, combined with the high melting point of RE-zirconates,  suggests  that RE additions may  be  a  useful  approach  to  improving the oxidation resistance of UHTC’s at  intermedi ate  temperatures  in  hypersonic  air.  In  terms  of  weight  change, the greatest improvement in the oxidation resistance of ZS20 at 1600°C reported herein was associated with the  use of Gd2O3 additions.  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},{
  "_id": 101,
  "PDF": "In situ microscopy observation of liquid flow, zirconia growth, and CO bubble formation during high temperature oxidation of zirconium diboride–silicon carbide.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 2365-2374  In situ microscopy observation of liquid flow, zirconia growth, and CO bubble formation during high temperature oxidation of zirconium diboride-silicon carbide  Sindhura Gangireddy a , Sigrun N. Karlsdottir b , S.J. Norton a , J.C. Tucker a , John W. Halloran a,∗  a Department of Materials Science and Engineering, University of Michigan, Ann Arbor, MI 48109, USA b Department of Materials, Biotechnology and Energy, Innovation Center Iceland, IS-112 Reykjavik, Iceland  Available online 20 February 2010  Abstract     The oxidation of ZrB2 -SiC composites at 1450-1650 C was directly observed with in situ optical microscopy. Video frames showed the flow of silicate liquids, the formation of zirconia deposits, and the growth and collapse of gaseous bubbles on the oxide surface. Contrast in the incandescence of in situ images is analyzed as spatial variations in hue and intensity and related to differences in emissivity of the oxide scale surface features by comparing these hot images with room temperature images. Above 1450 C, gaseous bubbles were observed to grow and collapse causing perturbations in the liquid oxide on the surface. The bubbles are associated with the evolution of CO from SiC oxidation and the onset is related to the critical temperature where the partial pressure of CO under the oxide scale exceeds atmospheric pressure. © 2010 Elsevier Ltd. All rights reserved.     Keywords: Optical microscopy; Borides; Refractories; Composites; In situ  1.  Introduction  1.1. Ribbon method  Ultra-high temperature ceramics (UHTCs) have recently gained interest as potential materials for a reusable thermal protection system and other components in hypersonic vehicles.1,2 Typically oxidation experiments have been conducted inside furnaces where it is not practical to directly observe the material while it undergoes high temperature oxidation. The other technique used in studying the high temperature oxidation behavior, arc-jet testing, simulates reentry environment and can allow viewing of the specimen. However, arc-jet testing is very expensive, not easily accessible and not typically instrumented for in situ studies.3 Consequently, oxidation processes have been inferred from post-test analysis of the oxide scale quenched to room temperature, regardless of the oxidation method.4,5 On the other hand, this study attempts to directly observe the UHTC specimen during high temperature oxidation using the Ribbon  ∗  Corresponding author. Tel.: +1 7347631051; fax: +1 734763478.  E-mail address: peterjon@umich.edu (J.W. Halloran).  0955-2219/$ - see front matter © 2010 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2010.01.034     Method.6 It has been verified that the ribbon method reproduces the complex oxide scales which form during high-temperature testing of UHTC. Owing to the special geometry of ribbon specimens and design of the process, only about 100 W of power is required to reach 1500-2000 C. As a result the heat flux from the sample is also small, enabling an optical microscope to approach the hot ribbon and image the surface at magnifications as high as 100×. An added advantage of ribbon method is that the fast heating rate (450 C/min) and free cooling rates (700 C/s)6 allow for samples to be tested easily in cyclic heating conditions while minimizing the effects of pre-oxidation caused by slower heating methods. The video images taken during oxidation showed gradual microstructural changes in the oxide scale due to liquid oxide flow and sudden changes caused by gas bubbles. In situ observations are compared with room temperature observations made after successive heating cycles.        1.2. Liquid convection hypothesis of oxide scale formation on ZrB2-SiC  Our observations are interpreted based on the dynamic evolution of crystalline and liquid features in the oxide in terms of the convective flow mechanism proposed by Karlsdottir et  \\x0c', '2366  S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  al.7 When ZrB2-SiC ultra-high temperature ceramic composites are oxidized at high temperatures, they form a complex oxide scale consisting of zirconia skeleton, covered by amorphous oxides, which act as protection barriers against oxidation.8,9 An explanation for these features has been given by a mechanism proposed by Karlsdottir et al.,7,10-12 which involves formation of a crystalline zirconia layer with a liquid oxide solution containing Boria-Silica-Zirconia (BSZ liquid), by the oxidation reaction: ZrB2 + xSiC + 5 + 3x O2 2 ⇒ y[ZrO2 ]primary+{B2O3+xSiO2+(1 − y)ZrO2 }liq + xCO  (1)     For a 15 vol% SiC composite, the condensed oxide products were estimated using a calculated 1500 C phase diagram.13 The predicted composition was one-third crystalline zirconia (molar basis) formed by direct oxidation (primary zirconia) and about two-thirds of the liquid oxide with approximate composition 71 mol% B2O3 + 18 mol% SiO2 + 11 mol% ZrO2 , which is referred to as the BSZ (Boria-Silica-Zirconia) liquid. However, boria is volatile at these temperatures and as it evaporates, the remaining liquid becomes rich in silica. {B2O3 + xSiO2 + (1 − y)ZrO2 }liq => {B2O3 }gas + (1 − y)ZrO2 + x{SiO2 }liq  (2)  The solubility of zirconia in silica-rich melts is lower12 than in borosilicates, so crystalline zirconia (secondary zirconia) precipitates from the liquid oxide as B2O3 evaporates. This secondary zirconia forms deposits at the surface of the oxide scale. Thus above the primary zirconia layer, the oxide scale consists of deposits of secondary zirconia, some fluid BSZ liquid, and a more viscous silica-rich liquid depleted of zirconia and boron oxide. Much smaller micron-sized zirconia particles, designated “tertiary zirconia” also decorate the surface. These tertiary zirconia particles seem to be carried along with the flowing liquids, serving as markers for the flow patterns. The oxide surface, observed under the in situ microscope, showed the formation of secondary zirconia and the flow patterns of the liquid glasses, confirming the liquid convection hypothesis.  2. Experimental procedure  2.1. Material fabrication  The UHTC material, ZrB2 -15 vol% SiC was provided by The Institute of Science and Technology for Ceramics (CNRISTEC) in Faenza, Italy. Details of the properties and processing have been presented elsewhere.14 The fabrication of the selfsupporting specimens is performed at University of Michigan involving the cutting of the bulk material with a wire-Electrical 2.3 mm × 2.0 mm × 25 mm Discharge Machine (EDM) into bars and then reducing the bar thickness to 400-500 \\u242em in the center with a mechanical grinder (220 grit - diamond wheel) to make the thin hot zones.6  2.2. Oxidation testing with in situ microscopy     The ribbon method6 is used to heat the samples to the testing temperature. The special specimen geometry for this method is shown in Fig. 1B. Electrical current is supplied through the thick ends, which heats the thin central ribbon to high temperature, while the end sections remain cold. The specimen is held by its thick ends onto the conductive specimen holder, as shown in Fig. 1C, with spring pressure from alligator clips (insulated with thin alumina plates). AC current is supplied through current leads to the conductive specimen holder, controlled by the temperature signal from an optical pyrometer (MI-S140, Mikron Infrared, Santa Clara, CA, USA). Details of the power supply, current and temperature control, and performance characteristics were previously reported.6 The samples are tested in temperature profiles of static oxidation and cyclic oxidation (in steps of 15 or 30 min), at temperatures ranging from 1450 to 1700 C. The in situ microscopy was performed using a stereo optical microscope (Nikon SMZ 1000, Nikon Instruments Inc., Melville, NY, USA) that was 15-20 cm above the specimen, length for different magnifications (1-8×). The microscope is shown in Fig. 1A. The exact lens distance varies with the focal fitted with a digital camera (01-GO-03-CLR-10, Q-Imaging, Surrey, Canada), and is connected through a computer to capturing software (“QCapture”, Q-Imaging, Surrey, Canada). The program provides an interface to view the sample in live preview as well as to capture still images or video. The incandescent radiation emitted from the hot sample is too bright for good visibility and therefore to reduce the intensity, a reflective neutral density filter of 2% optical density (NT48-531, Tech-Spec, Edmund Optics, Barrington, NJ, USA) is fitted over the microscope lens. This filter attenuates the light from the entire spectral range evenly and has a transmission ratio of 1%. The resulting light allowed viewing of features larger than 10 \\u242em. A video composed of frames snapped every 10 s, which is the allowed lower limit for QCapture programming, is recorded by the camera. The risk of missing important events occurring between the frames was small due to the relatively slow and gradual growth of the oxide scale features at temperatures below 1650 C, including bubble formation and associated liquid oxide motion.     2.3. Room temperature microscopy with cyclic heating     In cyclic heating, the specimen is cooled to room temperature almost instantly (700 C/s) after each cycle and then images are taken using the same stereo optical microscope, with external light focused on the sample, giving reflection contrast. These images taken after subsequent cycles of small time periods represent an interesting study of statistics in the formation and growth of the features. The samples were also observed under metallographic optical microscope which can reach higher magnifications and allows study of the microstructure of the oxide scale surface in more detail. Scanning electron microscopy (Philips XL30, FEI, Oregon, NE, USA) was also performed on the samples, but only after oxidation testing was completed since SEM required metallic conducting layer coating after which the sample can no longer be oxidation tested.  \\x0c', 'S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  2367  Fig. 1.  (A) In situ optical microscopy apparatus showing a stereo optical microscope situated over the heating component of the ribbon apparatus (UHTC specimen  at room temperature); (B) detail of the UHTC ribbon sample; (C) detail of the heating component of Ribbon method.  3. Results and discussion  3.1. Contrast in incandescence of in situ microscopy  4  to εL × TL  During the high temperature oxidation, in situ microscopy uses the incandescent light emitted by the hot specimen, which produces images that are quite different from ordinary optical microscopy images using reflected or transmitted light. The contrast in incandescent light microscopy is recorded as spatial variations in the wavelength (related to hue, H) and intensity, I, of the incandescent light. This information can be derived from the Red-Green-Blue (RGB) data in each pixel recorded by the camera, and can be quantitatively analyzed to extract hue and intensity.15 Hue (or color) is related to the wavelength of the emitted light which in turn varies with temperature of the source. Intensity, which is the amount of light in the radiation, is proportional (emissivity and local temperature raised to the fourth power). So hue and intensity are related to the spatial variations in local emissivity (εL ) and local temperature (TL ) where all hue variations correspond to variations in TL , and intensity variations - given constant hue, correspond to variations in εL . In other words, if the contrast is due to a difference in local temperature, the light and the dark regions differ in both intensity and hue values. On the other hand, if the contrast is due to a difference in local emissivity, intensity will change while hue will remain constant across these regions. Hue and intensity values of dark and light regions in the in situ images were compared over several locations of each in situ image and further over several images for verification. This comparison showed a shift in intensity while hue remained constant. So the contrast is concluded to arise from the difference in local emissivities of the dark and light regions. The surface of the oxide scale usually consists of crystalline zirconia and a silica-rich borosilicate liquid. The difference in the typical emissivities of these materials supports the above  analysis. The emissivity of crystalline zirconia ranges from 0.62 to 0.7516 whereas amorphous silica has emissivity ranging from 0.80 to 0.88.17 This suggests that the regions in brighter contrast must be silica-rich liquid and the features in darker contrast, crystalline zirconia. The areas selected for the analysis are all in the hot zone and are less than 50 \\u242em apart. So temperature differences within such small distances in hot zone should be negligibly small3 and would not have contributed for the contrast.  3.2. In situ image - comparison with room temperature images     Fig. 2A is an in situ image of the oxide scale surface on ZrB2-15% SiC at 1550 C after a total of 105 min of cyclic oxidation (15 min steps). The image shows the oxide surface imaged with its incandescent illumination. Fig. 2B shows the same field of view after quenching to room temperature, imaged in reflected light. The comparison of the images was used to identify the dark contrast features in the in situ image as secondary zirconia “islands” while the regions in lighter gray contrast are interpreted as silicate-rich liquid oxide, confirming the hypothesis made from contrast analysis correct. Curved lines, faintly visible in the incandescent image are interpreted as “flow lines” in the liquid oxide, the visibility of which is aided by small tertiary zirconia crystals entrained in the flowing liquid. The room temperature image shows the liquid oxide, now quenched to a glass, more clearly and with features suggesting that the liquid oxide emerged from near the zirconia islands and flowed across the surface. Fig. 3 compares two room temperature images taken of the hot zone showing the oxide scale surface on ZrB2 -15% SiC after cyclic oxidation in air at 1550 C for 60 min. The first is a reflected light optical image and the other is an SEM image. Zirconia is in bright contrast in the SEM image with silicate glass     \\x0c', '2368  S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  Fig. 2.  (A) In situ image at 1550     C after 105 min of oxidation, showing the oxide surface imaged with its incandescent  illumination. (B) The same field of view  after quenching to room temperature, imaged in reflected light.  in grey contrast. Both images showed clear distinction between turbid glass and clear glass. The liquid areas which are cloudy white in the optical image (and darker grey in the SEM image) have been shown7 to contain B2O3 , while the clear glassy areas in the optical image (and lighter grey in the SEM image) are richer in SiO2 . The appearance/location of these glassy regions suggests liquid flow outwards from the zirconia locations, closer to which the glass retained some boria (darker in SEM) whereas the glass farther from the zirconia was richer in SiO2 (lighter in SEM) owing to boria evaporation. Further evidence of flow comes from the patterns of micron-sized zirconia particles (tertiary zirconia) that decorate the surface, serving as markers for flow.  3.3. Dynamic evolution of oxide scale - in situ images     Fig. 4 shows a series of in situ images of the formation and growth of secondary zirconia islands taken during cyclic oxidation at 1550 C in incandescent illumination. After 1 min of oxidation (Fig. 4A) small secondary zirconia islands appear in dark contrast at locations (a), (b), (c) and (d). After 100 min of oxidation, Fig. 4B shows that these features have grown to about 50 \\u242em across. After 150 min oxidation (Fig. 4C), features (c) and (d) are growing together. Island size data was     collected from image frames for several specimens to determine the average diameter of secondary zirconia islands. The size of these features increases gradually with oxidation time up to about 180 min at 1550 C. Fig. 5 plots the average size from 50 ± 10 \\u242em after 15 min to 190 ± 10 \\u242em after 180 min. as a function of time and shows that the average size increases tions of 60 ± 10 \\u242em after 30 min and 115 ± 20 \\u242em at 180 min These values were comparable to previously reported observaon quenched samples for conventional furnace oxidation.11 So until 180 min, the average size increases roughly linearly with time at about 0.8-0.9 \\u242em/min. However, after the next cycle, the boundaries between the islands were less distinct and the average size increased from 190 to 230 \\u242em, a larger increase than expected from earlier cycles. This is associated with a sudden onset of bubbles forming during oxidation, which are presented in detail in the next section.  3.4. Observation of bubbles  The in situ video images taken during oxidation revealed formation and collapse of gas bubbles, called Bubble Burst Events, which had a dramatic effect on the oxide scale features. A bubble burst event, spanning from inception of a bubble until it ruptures, usually took about a minute to complete, owing to the high vis Fig. 3. Features on the surface of ZrB2 -15 SiC after conventional furnace oxidation in air at 1550 optical metallograph using reflected light (right).     C for 60 min. The same field of view imaged in SEM (left) and  \\x0c', 'S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  2369  Fig. 5. Diameter of secondary zirconia islands vs. oxidation time at 1550     C.     cosity of the liquid oxide film. Fig. 6 demonstrates the progress of a large bubble in the hot zone from the time of its formation, at 182 min of heating the specimen at 1550 C, until it burst one minute later. The sequence of images captures the slow growth of the bubble over the minute, demonstrated by increasing diameter of the bubble marked in dark contrast by the incandescence of in situ images. The three dark circles on the right hand side of these images are secondary zirconia. Fig. 6A is taken just before a bubble emerged in the lower left corner of the frame. Fig. 6B is a frame taken 20 s later, and the bubble about 250 \\u242em in diameter is seen in the lower left corner of the image. Fig. 6C is a frame another 40 s, and the bubble has grown to about 400 \\u242em diameter, and impinged upon the two secondary zirconia islands, causing the liquid oxide to flow over the zirconia. The bubble burst and collapsed shortly after this frame. The progress of the bubble boundary is outlined in Fig. 6D, with the position of the zirconia islands also noted. Fig. 7 shows the same location over the next few minutes of heating illustrating the motion of the liquid oxide caused by the above bubble burst event. This slow sweep of the liquid oxide from the vicinity of the bubble burst event, towards right side of the picture, lasted over 5 min. Fig. 7A is the image taken just after the bubble had burst, at 183 min of heating at 1550 C. Fig. 7B-D shows the same area 2 min later (185 min), 3 min later (186 min) and 5 min later (187 min) respectively. The shape of the flow pattern, marked (a), which is lined with tertiary zirconia particles in dark contrast, can be compared over this time duration. Fig. 7B shows a definite shift in (a) towards right side, the shift becoming pronounced with progress of time in Fig. 7C and D, as the liquid oxide keeps moving away from bubble burst location, pushing the flow pattern and submerging two secondary zirconia islands in the proximity, marked (b) and (c). The complete effect of this bubble burst is captured by two room temperature images (optical bright field) taken before, Fig. 7E and after the event, Fig. 7F. The two secondary zirconia islands are in the lower left, features (b) and (c), in Fig. 7E. However in Fig. 7F, (c) is almost covered by glass, and the island is not visible at location (b), as it has been covered by liquid oxide. In bright field, the silica-rich borosilicate glass reflects light off and only the features above the surface are visible. The glass     Fig. 4.  In situ images of the formation and growth of secondary zirconia islands     taken during cyclic oxidation at 1550  C after 1 min of oxidation (A), after  100 min (B), after 150 min (C). The regions of secondary zirconia appear  in  dark contrast at (a), (b), (c) and (d).  \\x0c', '2370  S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  Fig. 6.  In situ video frames of the formation and collapse of a gas bubble at 1550     C (A) after 182 min of oxidation; (B) after 182 min and 20 s; (C) after 182 min  and 60 s; (D) schematic showing the growth of the bubble with time until it burst and leaving the three zirconia deposits (represented as gray circles) covered with  liquid oxide.  that is now covering (b) and (c) has flow patterns with tertiary zirconia particulates, shifted from left due to the bubble burst event. Also, the zirconia island cluster on the right side seems to be larger. Thus bubble formation causes major dynamic changes in the oxide scale microstructure.  3.5. Trends in bubble burst events     The onset temperature and the number of bubbles were observed to depend on the temperature of oxidation. Constant heating is preferred over cyclic heating in studying these trends, owing to the brief spike in temperature by the PID controller while starting a heating profile which could influence the results which are being studied for temperature dependence. Bubbles were not observed below 1450 C even after extensive heating over 8 h, but appeared at 1500 C and above, suggesting an “onset temperature” for bubble formation around 1450-1500 C. It was also noted that bubbles do not appear immediately, but with a “delay time” requiring certain time of heating that occurs before bubbling is observed. This “delay time” was found to depend on temperature, shown in Fig. 8 (the dashed lines are to demonstrate the trend and are not most fit lines). While below 1450 they formed instantly, delay time  0. C no bubbles formed, above 1650 C The bubbles are usually accompanied with rapid oxidation rates and the sample burns through the thin cross-section of ribbon quite fast, leaving a short period of heating around              15-30 min when bubble bursts are occurring. During this period of time, the number of bubbles observed in the field of view is counted and normalized for “frequency of bubble formation”, expressed as the number of bubbles/cm2 s. In the temperature range where bubbles were observed, the frequency was found to increase steadily with temperature, as shown in Fig. 8. At 1450 C, where no bubbles were detected the frequency is zero and at 1650 C, the specimen has quite vigorous bubbling at a rate 16 bubble/cm2 s where the in situ videos showed the surface looking like boiling syrup.        3.6. Thermodynamic analysis of bubble formation  Let us consider the necessary conditions for forming a gas bubble. There must be a source of gas, either generated by a reaction or dissolved in a condensed fluid. To accumulate the gas, the rate of generation of the gas species must be larger than the rate it permeates out. The partial pressure of the species must be sufficient for a bubble to expand against the pressure of the atmosphere, so the pressure inside the bubble must be greater than Patm the external pressure (1 atm in this case). The possible gas species are CO, CO2 , SiO (from SiC) or a boron oxide gas (from ZrB2 ). Fahrenholtz estimated the vapor pressures of all boron oxide species to be much smaller than 1 atm at 1527 C.18 At the oxygen activity expected near ZrB2 /ZrO2 oxidation front 10 −16 ),18 the pressure of PCO2 will be much at 1527 less than PCO . At the same oxygen activity and temperature,  C (PO2        \\x0c', 'S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  2371  Fig. 7. A-D are in situ images taken during oxidation at 1550     C showing the oxide scale surface over the next 5 min after the bubble burst depicted in Fig. 6,  illustrating the motion of the liquid oxide away from the burst location. E and F are bright field optical images taken at room temperature before and after the bubble  burst events.  PSiO had been estimated to be <1 atm.19 Hence we suggest the gas bubbles are most likely composed of CO gas from the oxidation of silicon carbide. As SiC in the composite oxidizes to CO, one mole of CO is produced for every mole of SiC consumed. This CO can either be in the gas phase and be released as a bubble, or be dissolved in the condensed phases such as the BSZ liquid and permeate to the exterior by diffusion. Permeation of dissolved CO must be the common behavior, since bubbles are not typically observed.2,20 This may be because oxygen and CO might have similar permeabilities in oxide scales, owing to similar molecular diameters,21 and the oxidation of ZrB2-15% SiC requires the ingress of 2.725 mol of O2 but the egress of only 0.15 mol of CO. However, under certain conditions, the CO forms gas bubbles. Assuming that the activity of oxygen is same for both the  oxidation reactions, those of ZrB2 and SiC, we can estimate the partial pressure of CO in the bubble, PCO , by considering the equilibrium of ZrB2 with ZrO2 and B2O3 :  ZrB2 (s) + (5/2)O2 (g) = ZrO2 (s) + B2O3 (l)  With the equilibrium constant is given by  KZrB2  = aZrO2 aZrB2  = exp  aB2O3 P 5/2 O2  (cid:2) −\\x01G   f B2O3  (cid:3)  − \\x01G   f ZrO2  + \\x01G   f ZrB2  RT  (3)  (4)  in terms of the activities of the components, and the free energies of formation. If we consider passive oxidation of the SiC, the  \\x0c', '2372  S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  Fig. 8. The delay time before bubbles start to appear and the frequency of bubbles after the delay, while heating at different temperatures in static oxidation profile.  reaction is: SiC (s) + (3/2)O2 (g) = SiO2 (l) + CO (g)  With the equilibrium constant is given by  KSiC = aSiO2 PCO aSiCP 3/2 O2  = exp  (5)  (cid:3)  (cid:2) −\\x01G   − \\x01G   + \\x01G   f SiO2  f CO  f SiC  RT  (6)  oxidized at the same oxygen activity, we can replace the oxygen partial pressure term in (6) with expression from (4) to obtain an expression of the CO partial pressure under conditions where both equilibrium (3) and (5) obtain:  (cid:4)  (cid:5)3/5  PCO =  aZrO2 aB2O3 aZrB2  −1  (aSiO2 )  KSiC (KZrB2 )3/5  (7)  which involves both the oxygen activity and the CO partial pressure. If both the silicon carbide and the zirconium diboride are  in (5) are all 1, If we assume that the activity terms for the condensed species the PCO can be estimated simply from the  Fig. 9. Partial pressure of CO gas species, formed from passive oxidation of SiC, at the oxidation location as a function of temperature.  \\x0c', 'S. Gangireddy et al. / Journal of the European Ceramic Society 30 (2010) 2365-2374  standard free energies of formation:  (cid:2)  PCO =  KSiC (KZrB2 )3/5  = exp  \\x01G   f SiC  − \\x01G   f SiO2  − \\x01G   f CO  + (3/5)\\x01G   f B2O3  + (3/5)\\x01G   f ZrO2  − (3/5)\\x01G   f ZrB2  RT  (cid:3)  2373  (8)     The PCO thus calculated using data from the JANAF thermochemical tables22 is displayed in Fig. 9. Below 1450 C, the PCO is lower than the ambient pressure of 1 atm, but above 1450 C the PCO exceeds the minimum criterion for bubble to be stable. This transition temperature of 1450 C corresponds well with the experimentally observed “onset temperature” for bubble formation. Thus the thermodynamic analysis gives a justification for a temperature where bubbles begin, but does not seem to explain the time dependence or the “delay time” of Fig. 8.        3.7. Bubble burst events - cause or result?  It is not clear if CO bubbles increase the oxidation rate, or if they are a consequence of faster oxidation, but the bubbles seem to interrupt the slow growth process that dominates the first part of the oxidation and are associated with more rapid growth of the surface features, suggesting faster oxidation. One explanation could be that due to the large size of these bubbles, they can redistribute the liquid oxide and disrupt the protective oxide layer, leading to faster oxidation. Some of the samples show evidence of extensive bubbling and resulting retreat of the liquid oxide with exposed un-reacted material underneath. On the other hand, some other phenomenon might be causing an increase in oxidation rate, which implies increased rate of CO generation beyond the rate which can be removed by permeation and hence causing the bubbles. Nevertheless, the question remains why CO bubbles have not been described before. One possibility is that they are an artifact of the Ribbon Method. Unlike conventional furnace tests the Ribbon Method uses a specimen of limited thickness. It could be a phenomenon that occurs only in nearly fully oxidized ribbons. This method also uses direct resistive heating, where the interior of the sample is hotter than the surface, which could possibly cause bubbles that do not form in isothermal samples. However, it is also possible that the bubbles simply have not been noticed, or have not caught the interest of previous researchers. Perhaps the hot bulk ZrB2-SiC specimens have bubbles during conventhe 1 min lifetime of bubbles is tional furnace oxidation, but so brief that they collapse and disappear during cooling from furnace oxidation temperature.  4. Conclusion  Direct observations of high temperature oxidation behavior of ZrB2-SiC UHTC were achieved using in situ microscopy through a simple apparatus. Local variation in the incandescent light emitted from the hot sample provided the contrast in the images from the in situ optical microscopy. This contrast was analyzed to discover the origin to from the differences in the emissivity of crystalline zirconium dioxide and liquid silicon dioxide. In situ videos captured the evolution of the oxide scale ZrB2-SiC ultra-high temperature ceramics undergoing        oxidation, and showed the gradual growth zirconia regions and the flow of liquid silicate oxide, thus providing evidence of dynamic evolution of crystalline and liquid features in the oxide in terms of fluid flow mechanisms. The formation, growth, and collapse of bubbles during the oxidation of ZrB2-15% SiC were at 1450-1650 C were observed. No bubbles were reported below an onset temperature of 1450 C. Above this temperature, bubbles were not observed until after a delay time, an extended period of oxidation. The delay time decreased at higher temperatures and the frequency of bubble formation was found to increase with temperature of oxidation. Thermodynamic calculation of the carbon monoxide partial pressure presumed to exist at the boride + carbide/oxide interface showed that the CO pressure should exceed one atmosphere at around 1450 C, which is close to the experimentally observed onset temperature. Overall, the emergence, growth, and collapse of the bubbles cause significant motion of the liquid oxide, and drastic changes in the scale microstructure.     Acknowledgement  We thank the Office of Naval Research for research under grant N00014-02-1-0034.  funding this  References  1. Zimmermann JW, Hilmas GE, Fahrenholtz WG, Monteverde F, Bellosi A.  Fabrication and properties of reactively hot pressed ZrB2 -SiC ceramics. J Eur Ceram Soc 2007;27(7):2729-36.  2. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during the  oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27(6):2495-501.  3. Karlsdottir SN. 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Molecular  theory of gases and  ceramics and composites. 2007. p. 327-36.  liquids. New York: John Wiley & Sons Inc.; 1954. p. 165, 599.  15. Russ JC. The image processing handbook. second ed. Florida: CRC Press;  22. Stull DR, Prophet H, NSRDS, JANAF Thermochemical Tables, National  1995. p. 33-45.  Bureau of Standards, NBS 27, second ed., Washington DC, 1971.  \\x0c']"
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  "_id": 102,
  "PDF": "In situ studies of oxidation of ZrB2 and ZrB2-SiC composites at high temperatures.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 2375-2386  In situ studies of oxidation of ZrB2 and ZrB2-SiC composites at high temperatures  P. Sarin, P.E. Driemeyer, R.P. Haggerty, D.-K. Kim, J.L. Bell, Z.D. Apostolov, W.M. Kriven  Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, 1304 West Green Street, Urbana, IL 61801, USA  Available online 15 April 2010  Abstract  ∗     High temperature oxidation of ZrB2 and the effect of SiC on controlling the oxidation of ZrB2 in ZrB2 -SiC composites were studied in situ, in air, using X-ray diffraction. Oxidation was studied by quantitatively analyzing the crystalline phase changes in the samples, both non-isothermally, temperature, up to 1650 time, at 1300 as a function of C, as well as isothermally, as a function of C. During the non-isothermal studies, the formation and transformation of intermediate crystalline phases of ZrO2 were also observed. The change in SiC content, during isothermal oxidation studies of ZrB2 -SiC composites, was similar in the examined temperature range, regardless of sample microstructure and composition. Higher SiC content, however, markedly retarded the oxidation rate of the ZrB2 phase in the composites. A novel approach to quantify the extent of oxidation by estimating the thickness of the oxidation layer formed during oxidation of ZrB2 and ZrB2 -SiC composites, based on fractional conversion of ZrB2 to ZrO2 in situ, is presented. © 2010 Elsevier Ltd. All rights reserved.     Keywords: Ultra-high temperature ceramics; X-ray methods; Borides; ZrB2 ; SiC; Oxidation  1.  Introduction  Transition metal diborides have been subjects of extensive research in the last decade as ultra-high temperature ceramics (UHTC’s) for applications such as thermal protection systems on re-usable atmospheric re-entry vehicles and on scramjet engine components for hypersonic aerospace vehicles.1-6 These materials exhibit a unique combination of thermo-mechanical properties which make them attractive candidates for aerospace applications under extreme environments. Zirconium diboride (ZrB2 ) is one such UHTC which has been studied because of its extremely high melting temperature (Tmelt > 3040 C), high elastic modulus, high electrical and thermal conductivity, good thermal shock and wear resistance, and chemical inertness.7-11 The superior high temperature properties are primarily due to strong covalent bonding in the ZrB2 structure.11 Extremely low self-diffusion coefficients, on the other hand, make processing of dense ZrB2 ceramics a challenge. Pure ZrB2 is typically sintered around 2000 C under an applied load in a controlled atmosphere furnace.12 Numerous studies have focused on the use of alternate sintering methods to achieve dense ZrB2 ceramics, such as        ∗  Corresponding author. Tel.: +1 217 333 5258; fax: +1 217 333 2736.  E-mail address: kriven@illinois.edu (W.M. Kriven).  0955-2219/$ - see front matter © 2010 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2010.03.009  by using sintering aids, by reactive routes or by spark plasma sintering.11 Another major challenge in the development of ZrB2 UHTCs lies in their poor resistance to oxidation which severely undermines their ability to survive oxidizing conditions at high temperatures. When exposed to air, crystalline ZrB2 oxidizes to ZrO2 and B2O3 according to reaction (1).13 At temperatures less than 1100 C, the B2O3 forms a continuous liquid layer which limits the transport of oxygen to the ZrB2 surface, and as a consequence diffusion controlled, parabolic, oxidation kinetics are observed.14-19 The ZrO2 formed on oxidation of ZrB2 is porous and therefore non-protective towards further oxidation:     ZrB2 + 5 2 O2 (g) → ZrO2 + B2O3 (l)  (1)        Based on thermo-gravimetric analysis (TGA), at temperatures above 1100 C, the oxidation rate increases and exhibits para-linear kinetics between 1100 and 1400 C. In this temperature range, the mass change is determined by (a) weight gain due to the formation of B2O3 and ZrO2 and (b) weight loss due to vaporization of B2O3 (l).17,18 At temperatures higher than 1400 C, linear oxidation kinetics ensue due to rapid evaporation of B2O3 , and a non-protective porous ZrO2 scale is formed.20,21 Overall, a net gain in mass is observed.     \\x0c', '2376  P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386           The addition of silica scale formers, such as SiC or MoSi2 , has been shown to improve oxidation resistance above 1100 C.19,22-27 The inclusion of SiC results in the formation of a silica (SiO2 ) based glassy layer on the ZrB2-SiC composite surfaces. The SiO2 layer is less volatile than B2O3 above 1100 C, resulting in slow, diffusion-controlled oxidation kinetics in ZrB2-SiC composites over a much greater temperature range.23 Moreover, SiC acts as a grain growth inhibitor, increasing sinterability, and thus improving strength, toughness, and oxidation resistance of the composite.6,8,17,22,23,28,29 Oxidation of ZrB2-SiC at 1500 C produces a multilayered microstructure consisting of a surface scale that is SiO2 rich, followed by a layer of non-oxidized ZrB2 over the underlying ZrB2 -SiC composite. In addition, a layer composed of crystalline ZrO2 has also been reported between the silica-rich surface layer and the non-oxidized ZrB2 layer.6,30,31 The formation of such a layered structure exhibiting a “SiC-depleted” region above the ZrB2-SiC bulk composite during oxidation has been noted by several authors.19,20,23,24,30 The absence of SiO2 or other condensed phases in the SiC-depleted zone has been explained by its removal as gaseous phases such as SiO.30 Most of the current understanding on the oxidation of ZrB2 and ZrB2-SiC composites is based on thermo-gravimetric studies, oxygen consumption studies and on ex situ evaluation of microstructural changes. While thermo-gravimetric methods provide useful information based on weight changes during oxidation, they record the simultaneous occurrence of multiple events including weight gain due to formation of oxide phases, such as ZrO2 , B2O3 and SiO2 , and weight loss due to vaporization of B2O3 and perhaps SiO. As a result, the TGA data needs to be deconvoluted and requires careful interpretation. The post-oxidation, ex situ analysis of microstructure, on the other hand, is largely qualitative and limited in knowledge of the oxidation phenomena and reactions, as they occur. The oxidation and the scale formation in the UHTC diborides involve the complex interaction of a number of factors. The resulting multiphase, multilayered scale microstructure makes it even more difficult to determine the factors that control oxygen transport during the oxidation of the composites. In order to realize the full potential of UHTC diborides for advanced hypersonic and propulsion applications, and to improve their oxidation resistance, a more comprehensive understanding of the oxidation process is desired. In situ high temperature investigation of the oxidation of UHTC diborides using X-ray diffraction can provide a clearer insight into the oxidation kinetics and structure-property relationships. Although such studies are extremely challenging due     to instrumental limitations, they can lead to improved life prediction of UHTC components and perhaps indicate an optimal composite through tailoring of composition and microstructure. This study was aimed at evaluating the use of high temperature X-ray diffraction (HTXRD) to examine the oxidation of ZrB2 up to 1650 C in air. The effect of SiC on the oxidation of ZrB2 in ZrB2-SiC composites was also evaluated. A specially designed quadrupole lamp furnace (QLF), which allows heating of samples up to 2000 C in air,32-35 was used in conjunction with a high resolution curved image plate (CIP) detector for these studies.36,37 Using this method, the oxidation of ZrB2 to ZrO2 was monitored in real time, independent of other accompanying processes.     2. Experimental procedure  2.1.  Sample processing  Commercially available ZrB2 (Aldrich Chemical Company, Milwaukee, WI, USA) with a reported purity of 95% (purity excludes 1-2% Hf), density of 6.08 g/cm3 , and an averaged particle size of 5 \\u242em was used. The SiC powder (Aldrich Chemical Company, Milwaukee, WI, USA) used in this study was predominantly ␤-SiC (3-15% amorphous), had a density of 3.22 g/cm3 , an average particle size <0.1 \\u242em and specific surface area of 70-90 m2 /g. While the pure ZrB2 sample was made using the commercially available powder, two ZrB2-SiC composites were prepared by mixing batches of commercial powders of ZrB2 and SiC in proportion such that the final samples had ZrB2 and SiC in the ratios 70:30 and 50:50 by volume, respectively. The powder batches for the composite samples were ball milled for 24 h, using 3Y-TZP (3 mol% Y2O3 stabilized tetragonal zirconia (ZrO2 ) polycrystalline) balls of 5 mm diameter as the milling media, to reduce particle size and achieve homogenous mixing. The powders of pure ZrB2 phase and the ZrB2-SiC composite samples were first uniaxially pressed into disks with an approximate diameter of 25 mm under a pressure of 35 MPa. The pressed disks were then subjected to cold isostatic pressing at 414 MPa for further compaction. The samples were then hot pressed according to conditions detailed in Table 1. While the ZrB2 samples were hot pressed at 1700 C for 3 h under 35 MPa of pressure, the ZrB2-SiC composite samples were hot pressed at 1900 C for 2 h under 35 MPa pressure. The entire hot pressing was performed in an Ar atmosphere and a heating rate of 50 C/min was used. The furnace was turned off after holding for the specified time at the targeted temperature. Disks with a diameter of 25 mm and thickness of           Table 1  Raw materials and processing conditions for preparation of ZrB2 and ZrB2 -SiC composite samples.  Sample  Composition (vol. %)  ZrB2 70:30 ZrB2 -SiC 50:50 ZrB2 -SiC  ZrB2  100  70  50  SiC  0  30  50  a 3 mol% Y2O3 stabilized tetragonal zirconia (ZrO2 ) polycrystals.  Powder preparation  Ball milling  None 24h; 3Y-TZPa media 24h; 3Y-TZPa media  Drying       -  C  C  150  150  Hot pressing  Temp./time/pressure         1700  1900  1900  C/3 h/35 MPa  C/2 h/35 MPa  C/2 h/35 MPa  \\x0c', 'P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  2377  5 mm were produced. Smaller specimens having dimensions of 5 mm × 3 mm × 0.5 mm were cut from the disks by using a slow action diamond saw (Isomet, Buehler Ltd., Evanston, IL, USA) and polished to a 6 \\u242em surface finish. The final samples for microstructure characterization and oxidation studies using TGA, differential scanning calorimetry (DSC), and HTXRD, had dimensions of 5 mm × 3 mm × 0.3 mm.  2.2. Microstructure characterization     The physico-chemical characteristics of the hot pressed samples were determined using a variety of methods. Bulk density, apparent specific gravity and apparent porosity of the hot pressed samples were measured by the Archimedes method, following the procedure detailed in ASTM C20.38 The samples used for these measurements were semi-circular in shape (diameter 25 mm, height 5 mm) and were prepared by sectioning the hot pressed sample disks using the slow action diamond saw, followed by ultrasonic cleaning. All samples were dried at 105 C for 24 h in air before any measurements were taken. Although the sample sizes used for these measurements were smaller than those recommended in ASTM C20, they were considered adequate for the purpose of this study. Distilled water was used as the suspending medium. The microstructure and elemental composition of hot pressed and oxidized samples were characterized by scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) using a Hitachi S-4700 (Hitachi High Technologies, Schaumburg, IL, USA) and a JEOL JSM-6060LV (JEOL USA, Inc., Peabody, MA, USA) scanning electron microscopes (SEMs). The hot pressed samples of ZrB2 and ZrB2-SiC composites used for SEM/EDS analyses had dimensions of 5 mm × 3 mm × 0.3 mm, and were polished to a 1 \\u242em surface finish on Buehler’s ECOMET III polisher/grinder (Buehler Ltd., Evanston, IL, USA) using a diamond paste (Buehler METADI Aerosol Spray Diamond Compound—1 Micron; Buehler Ltd., Lake Bluff, IL, USA). The oxidized ZrB2 and ZrB2-SiC composite samples were first impregnated with an epoxy resin (Epo-Thin, Buehler Ltd., Lake Bluff, IL, USA), and then sectioned to enable examination of the cross-section. These samples were also polished to a 1 \\u242em surface finish using the diamond paste. Water was used as a lubricant during sectioning and polishing. All SEM samples were mounted on aluminum stubs and sputter coated with 6 nm of a Au/Pd alloy to facilitate imaging. The samples used for EDS analysis were coated with carbon instead of Au/Pd, and flat sample regions were examined using a 20 kV or 15 kV accelerating voltage and a 10 mm working distance. Copper was used to calibrate the energy and a minimum of six acquisitions were taken on each sample. The crystalline composition of the hot pressed and oxidized samples was determined by synchrotron X-ray diffraction (XRD). Plate-shaped samples (with dimensions of 5 mm × 3 mm × 0.3 mm) were studied in transmission geometry. Further details on the instrumentation and calibration of the XRD experiments are included in Section 2.3.2. The XRD patterns were subsequently analyzed by the Rietveld method39 using the JADE software (Materials Data, Inc., Livermore, CA,  USA) to determine the crystalline phase composition as well as crystallographic parameters for each constituent phase.  2.3. Oxidation studies  The oxidation of ZrB2 and ZrB2-SiC composites was studstudies were hot pressed plates (5 mm × 3 mm × 0.3 mm) which ied using TGA/DSC and HTXRD. The samples used for these had been polished to a 6 \\u242em surface finish.  2.3.1. TGA/DSC studies     Simultaneous TGA and DSC studies were conducted on hot pressed ZrB2 and ZrB2-SiC composites. Samples were heated at 10 C/min in a Netzsch DSC/TGA (Model STA409 CDTM , Export, PA, USA) instrument. While the ZrB2 sample was heated up to 1350 C, the ZrB2-SiC composite samples were tested up to 1500 C. An alumina (Al2O3 ) pan fitted with a lid was used to hold the specimen and as a reference. During the analysis, the sample chamber was purged with He (25 mL/min) and air (50 mL/min).        2.3.2. In situ high temperature synchrotron diffraction studies  In a typical HTXRD experiment, the hot pressed plate specimens were heated, and the crystalline composition and structural changes were simultaneously recorded as XRD patterns using synchrotron radiation and a curved image plate (CIP) detector.36 All experiments were conducted at the 33BM-C beam line at the Advanced Photon Source (APS) at Argonne National Laboratory, Argonne, IL, USA. The experimental set-up used for these studies is shown in Fig. 1. Plate samples were mounted vertically, perpendicular to the incident X-ray beam, in a specially designed sample holder made with Pt-20%Rh and heated using a quadrupole lamp furnace (QLF).35 The schematic of the sample holder along with the specimen is shown in Fig. 2. The sample was coated with Pt powder (0.15-0.45 \\u242em, 99.999% pure, Aldrich, Milwaukee, WI, USA) on the side facing the incident X-ray beam and was used to determine sample temperature from the expansion of the Pt lattice. The QLF is a radiation heating furnace that can be used to conduct HTXRD investigations of ceramics in air up to 2000 C in air.32-35 The details of the construction, operation and capabilities of the CIP detector, used in these studies, are reported elsewhere.36 The CIP detector allows for simultaneous acquisition of diffracted X-ray intensities over a 2θ range extending from 2 to 35 , thus eliminating any time dependence in XRD pattern acquisition. Contingent upon sample properties and incident X-ray beam intensity, high resolution XRD patterns can be acquired in ≤20 s using this detector.37 For the purpose of this study, the CIP detector was first aligned and calibrated using a Si(1 1 1) analyzer crystal and the LaB6 powder standard (SRM 660a, National Institute of Standards and Technology, Gaithersburg, MD, USA). Incident monochromatic X-rays of wavelength 0.70087 Å, as calibrated with SRM 660a, were used in this work. Oxidation of ZrB2 was studied by heating the plate sample in steps of 100 in air to discreet set temperatures (Tset ) C, from room temperature up to 1500 C, and X-ray diffraction                 \\x0c', '2378  P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386              Tset  crystallites, the plate samples were rocked ±1 about a horizontal axis perpendicular to the incident X-ray beam. The sample was cooled down to room temperature after the highest temperC) in 5 min and XRD patterns were ature run (i.e. Tset = 1500 recorded to determine the final crystalline phase composition after oxidation. Isothermal HTXRD studies were conducted in air in order to evaluate the effect of SiC on the oxidation of ZrB2 . Plate samples of ZrB2 and ZrB2-SiC composites were heated to Tset = 1100 C in 5 minutes, and several HTXRD patterns were acquired over the next 1 h duration while the samples were maintained at the temperature. Based on prior experience and the location of the control thermocouple (which measures Tset ) in the QLF, the sample temperature (Tsample ) during these studies was expected to be in the range 1200-1350 C.32-35,37 This temperature was selected for the isothermal studies as it corresponded to oxidizing conditions in a partial protective regime for the materials being evaluated. Therefore, it was anticipated that a pronounced effect of the presence of SiC on oxidation of ZrB2 could be suitably observed. Following the isothermal studies, each sample was cooled down to room temperature in 5 min. XRD patterns were also collected at room temperatures for the ZrB2 and ZrB2-SiC composite samples, both before and after the isothermal studies, to determine the initial and final crystalline phase composition. All the XRD patterns, including those acquired at room temperatures and at high temperatures, were analyzed by the Rietveld method using the JADE software (Materials Data, Inc., Livermore, CA, USA) to extract the quantitative crystalline phase composition and crystallographic parameters. The PDF-4 2008 database from ICDD (International Center for Diffraction Data, Newtown Square, PA, USA) and the Inorganic Crystal Structure Database (NIST, Gaithersburg, MD, USA; and Fachinformationszentrum (FIZ), Karlsruhe, Germany) were used for crystalline phase identification. The structure parameters of the crystalline phases identified from the databases were used as starting parameters during refinement. Besides the lattice constants for each phase, other factors that were refined included scaling factors, sample displacement, profile function parameters, and isotropic temperature factors (where possible).  Fig. 1. Photograph of the HTXRD experiment set-up at the 33BM-C beam line  at APS, Argonne National Laboratory. A curved image plate (CIP) detector was  used with a quadrupole lamp furnace (QLF).  patterns were acquired in situ, in transmission geometry at each temperature step. At least four HTXRD patterns were collected and averaged at each temperature, to improve signal-to-noise ratio. The sample temperature (Tsample ), which is different from the set temperature (Tset ) when using the QLF,35 was determined later during analysis of HTXRD patterns from the lattice expansion of cubic Pt. In order to maximize random orientation of the  3. Results and discussion  3.1. Microstructure of processed samples  The microstructure of the hot pressed ZrB2 and the ZrB2-SiC composites was fairly porous as is evident from the SEM micrographs presented in Fig. 3. In the case of the 70:30 ZrB2 -SiC and the 50:50 ZrB2-SiC samples, SiC was fairly uniformly distributed in the ZrB2 matrix (see inset in Fig. 3(b) and (c)). The grain size of the ZrB2 phase was larger in the pure ZrB2 specimen (Fig. 3(a)) in comparison to the composite samples. This is most likely due to the presence of SiC which can act as a grain growth inhibitor. Moreover, the latter samples had also been subjected to 24 h of ball milling during sample processing. The crystalline phase composition of each of the samples was  Fig. 2. Schematic of the sample holder used to mount hot pressed plate samples  for HTXRD studies in transmission geometry.  \\x0c', 'P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  2379  Fig. 3. SEM micrographs of the processed samples: (a) ZrB2 , (b) 70:30 ZrB2 -SiC, and (c) 50:50 ZrB2 -SiC. Elemental maps of Zr and Si are shown as insets for 70:30 ZrB2 -SiC and 50:50 ZrB2 -SiC samples.  verified by Rietveld analysis of the XRD patterns collected at room temperature for each sample. The results are included in Table 2 and they confirmed that the hot pressed composite samples had the component phases in the desired proportion. There was no evidence of any impurities or of the formation of any oxidized crystalline phases during sample processing. The density and porosity of the hot pressed samples, measured using the Archimedes’ method, were consistent with the observed microstructures (see Table 2). While the pure ZrB2 and the 50:50 ZrB2-SiC samples had >30% porosity, the 70:30 ZrB2-SiC composite samples were 23.7% porous. The measured apparent densities of the composite samples were consistent with the expected values calculated using densities of each component phase and the phase composition. Therefore, it was concluded that porosity in the composite samples was predominantly interconnected. However, considerably low apparent density value of the pure ZrB2 sample, in comparison to pure ZrB2 phase, suggests the presence of closed pores. Since this study was aimed at examining the oxidation behavior of hot pressed ZrB2 and ZrB2-SiC, samples with interconnected  porosity were considered particularly suitable to measure oxidation kinetics in reasonable time due to the increased surface area available for oxidation. This argument is especially relevant for HTXRD studies using synchrotron radiation, as experimental time is limited.  3.2. Oxidation studies using TGA/DSC     Percentage weight changes vs. temperature of the ZrB2 and the ZrB2 -SiC composites during non-isothermal heating up to 1500 C are presented in Fig. 4(a). The heating rate of 10 C/min allowed for detection of suitable DSC signal to assess any exothermic or endothermic processes occurring during oxidation. No change in the samples was observed up to 650 C, which was followed by a rapid increase in mass due to oxidation. This was also detected as an exothermic event in the DSC signal. The onset of oxidation was delayed in the ZrB2-SiC samples, and Tstart of 661, 670, and 691 C were recorded for the pure ZrB2 , 70:30 ZrB2 -SiC and the 50:50 ZrB2-SiC samples, respectively. The delay in the onset of oxidation of ZrB2 in the           Table 2  Physical properties and crystalline phase composition of hot pressed ZrB2 and ZrB2 -SiC composite samples.  Sample  Density (g/cm3 )  Total porosity (vol. %)  Crystalline phase composition from XRD  ρbulk  ρapparent  ZrB2 70:30 ZrB2 -SiC 50:50 ZrB2 -SiC  3.84  4.07  2.96  5.68  5.34  4.63  32.3  23.7  35.9  Weight %  ZrB2  SiC  100 82.6 ± 0.7 67.4 ± 0.4  17.4 ± 0.4 0 32.6 ± 0.3  Volume %  ZrB2  100 71.5 ± 0.3 52.3 ± 0.1  SiC  28.5 ± 0.1 0 47.7 ± 0.1  \\x0c', '2380  P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  Fig. 4. Thermal analysis studies on ZrB2 and ZrB2 -SiC composite samples, from 20 to 1450     C in air: (a) TGA and (b) DSC.     ZrB2 -SiC composites is most likely due to coating of the ZrB2 grains with SiC, as seen in the SEM micrographs (Fig. 3(b) and (c)). The rapid increase in weight observed in all the samples between 700 and 800 C is consistent with the high porosity and larger surface area, which ensured the availability of the ZrB2 phase for active oxidation. The 50:50 ZrB2-SiC samples were more porous than the 70:30 ZrB2-SiC samples (see Table 2). The overall weight gain observed in the 50:50 ZrB2-SiC sample, in the studied temperature range, was larger than the 70:30 ZrB2-SiC sample. This could be a result of higher SiC content in the sample. It is reasonable that in the investigated temperature range the net gain in weight in the 50:50 ZrB2-SiC composite sample is primarily determined by the oxidation of SiC. This can be comprehended better by considering complete oxidation of ZrB2 to ZrO2 and SiC to SiO2 in the samples, along with the loss of B2O3 due to vaporization. With this assumption, the final increase in weight for the pure ZrB2 sample, the 70:30 and the 50:50 ZrB2-SiC composites is expected to be 9.2%, 16.6% and 22.3%, respectively. In this case, the weight gain in both the calculated as 9% (of the 16.6%) for the 70:30 ZrB2-SiC samcomposite samples is dominated by the oxidation of SiC, and is ple and 16% (of the 22.3%) for the 50:50 ZrB2-SiC sample. The loss of Si as SiO species is assumed to be negligible in this argument. Based on these TGA/DSC studies alone, it is difficult to judge the efficacy of SiC content in protecting ZrB2 from oxidation. The change in slopes of the TGA curves at around 800 C in the composite samples, and again at 900 C, suggest a slower rate of oxidation which is perhaps due to the transient oxidation protection afforded by the liquid B2O3 phase. Beyond 1200 C, the increase in slope of the TGA curves for the pure ZrB2 and the 70:30 ZrB2-SiC samples suggest rapid oxidation due to vaporization of B2O3 phase. On the other hand, the decrease in the rate of weight gain observed in the 50:50 ZrB2 sample is most likely due to formation of a glassy SiO2 layer which protects rapid oxidation of the underlying ZrB2 phase. The DSC data is unremarkable except for the large exothermic peaks due to oxidation reactions. The presence of a small endothermic peak at 1200 C, observed only in the case of the pure ZrB2 sample, does suggest a possible transformation of the monoclinic ZrO2              (m-ZrO2 ) phase, expected to be formed on oxidation of ZrB2 , to tetragonal ZrO2 (t-ZrO2 ) phase.40,41  3.3. Oxidation of ZrB2—crystalline phase evolution with temperature                    The evolution of crystalline phases during oxidation of pure ZrB2 sample in air was recorded as a series of HTXRD patterns in discreet Tset steps of 100 C, from room temperature to Tset = 1500 C in air. Although HTXRD datasets were collected over a 2θ range from 2 to 35 , only segments of the dataset are presented in Fig. 5 for clarity. Each HTXRD pattern presented in Fig. 5 was averaged over 4 different exposures which were recorded in ≤4 min total time, while the sample was maintained at the specified Tset . The set temperature was increased in steps of 100 C in 1 min, followed by equilibration for 1 min before acquisition of 4 different exposures using the CIP detector. Based on lattice expansion of the Pt phase, which was coated on one side of the sample, the oxidation behavior of ZrB2 was studied in the temperature range extending from 20 to 1627 C in air. Qualitative changes occurring during oxidation of ZrB2 are clearly evident from the HTXRD patterns in Fig. 5. The drift of the diffraction peaks for all the phases towards lower 2θ angles, is the result of temperature induced expansion of crystal lattices of each phase. The change in the Pt peak around 29 , starting as a broad peak at 20 C to a well defined sharp peak, is a result of Pt crystallite growth with increase in temperature. The first crystalline oxide phase as a result of ZrB2 oxidation was found to be t-ZrO2 (ICDD PDF#04-005-4504). Subsequent appearance and disappearance of the m-ZrO2 (ICDD PDF # 01-083-0936) peaks due to transformation into t-ZrO2 phase at higher temperatures C) was also observed. Although not apparent from the HTXRD patterns in Fig. 5, a broad peak, symbolic of the presence of an amorphous phase, was observed in patterns acquired at Tsample ≥ 935 C (or Tset ≥ 700 C). The analysis of the HTXRD patterns using the Rietveld method provides a quantitative insight into the oxidation of ZrB2 at high temperatures. The results are presented in Fig. 6 as weight percentage composition of crystalline phases observed  (Tsample ≥ 1250                 \\x0c', 'P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  2381  Fig. 5. High temperature XRD study of ZrB2 , from 20 to 1627     C, in air.     at each temperature. The smooth lines drawn through the data points are merely a guide to assist in clearer depiction of the results and should not be misconstrued to represent composition at intermediate temperatures between any two measured data points. No oxidation of ZrB2 was observed up to 800 ± 3 C, beyond which the ZrB2 content of the sample decreased continuously, resulting in the formation of t-ZrO2 or m-ZrO2 phases. As stated above, the first oxidized phase to be observed was t-ZrO2 which appeared as broad peaks, most notably around , at 906 ± 1 13.5 C. The overall composition at this temperature was 81.7 ± 0.5% ZrB2 (ICDD PDF # 04-004-7151) and 18.3 ± 0.3% t-ZrO2 by weight. Upon heating to 997 ± 1 C the m-ZrO2 phase was formed and its concentration increased from 13.1 ± 0.3% to 33.8 ± 0.3% (weight %) at 1169 ± 1 C, while the t-ZrO2 content decreased from 18.8 ± 0.2% to 4.2 ± 0.1% (weight %) in the same temperature range. Heating the sample to higher m-ZrO2 to t-ZrO2 by 1250 ± 1 temperatures resulted in complete conversion of C, and only t-ZrO2 phase was                             observed at higher temperatures. The ZrB2 phase in the sample was completely oxidized by 1477 ± 1 C. The observation of t-ZrO2 at temperatures below the transformation temperature (i.e. 1170 C) is most likely due to the particle size effect, which has been previously reported.40,41 A reasonable estimate of the average crystallite size, or coherently scattering crystalline domains, can be made from the broadening of diffraction peaks, which is attributable only to the sample. In the case of ceramic phases, crystallite sizes are often similar to grain sizes. The average crystallite size of culated to be 19 nm at 906 ± 1 t-ZrO2 was cal1169 ± 1 C, which increased to 73 nm at C. This estimation was made after taking into account the instrumental broadening, which was determined using SRM 660a LaB6 powder. Since this calculation made use of the Scherformula,42 any microstrains present were ignored, and the rer crystallite size values may be underestimated. Nonetheless, crystallite size measurements are consistent with the presence of t-ZrO2 phase at lower temperatures than expected. The presence of m-ZrO2 phase up to 1250 ± 1 C, even though only in small amounts (i.e. 3.6 ± 0.1% by weight), is intriguing. It is expected that all the m-ZrO2 phase transforms to t-ZrO2 phase by 1170 C.40,41 However, it is possible that due to impurities such as Hf in the starting ZrB2 phase, this anomaly was observed. It is also conceivable that due to slow diffusion of oxygen through a liquid B2O3 layer a defect m-ZrO2 structure is formed with oxygen vacancies resulting in its stability up to higher temperatures. Overall, non-isothermal HTXRD studies were quite similar to the TGA/DSC studies reported in Section 3.2, but provided clearer and quantitative insight into oxidation of ZrB2 , and the formation and transformations of the crystalline oxidized phases.        Fig. 6. Crystalline phase evolution with temperature during oxidation of ZrB2 in air.  3.4.  Isothermal oxidation studies  Isothermal oxidation studies of ZrB2 and ZrB2-SiC composites using HTXRD were conducted in order to (a) evaluate the usefulness of the method towards understanding oxidation kinetics of UHTC diborides by monitoring changes in the crystalline phase composition and (b) to examine the effect of SiC on  \\x0c', '2382  P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  Fig. 7.  Isothermal high temperature XRD study of ZrB2 at 1268 ± 3     C in air.              the oxidation kinetics of ZrB2 in ZrB2 -SiC composite ceramics. Once again, the sample temperature was determined from the lattice expansion of Pt. Fig. 7 shows the series of HTXRD patterns that were collected at 1268 ± 3 C to observe oxidation of pure ZrB2 . Each of the thirty-seven HTXRD patterns presented in Fig. 7 was recorded using the CIP detector in approximately 1 min 41 s, while the sample was maintained at the specified temperature. The first HTXRD pattern was collected after the sample had been heated at 1268 ± 3 C for 198 s, while the last HTXRD pattern at this temperature was recorded 3678 s after reaching the temperature. The sample was exposed to the X-ray beam for 30 seconds for each HTXRD pattern. Only segments of the dataset are presented, although HTXRD datasets were collected over a 2θ range from 2 to 35 . Each subsequent HTXRD pattern has been offset, both along the vertical direction as well as to the right along the 2θ axis, to clearly show the growing and the diminishing crystalline phase peaks belonging to t-ZrO2 and the ZrB2 phases, respectively. No change was observed in the Pt peak location or intensity during the entire experiment. The ZrB2 peak intensities gradually decreased with time of exposure to high temperature, while the t-ZrO2 phase peaks consistently increased in intensity. Occasionally, a few peaks for ZrB2 phase defied the trend and have higher intensity than observed in the preceding patterns, for example the peaks near 14.5 , 18.5 and 22.6 . This is most likely due to transient preferred orientation recorded during the 30 s exposure to X-ray beam, and is due to the inability to achieve completely random orientation of the diffracting ZrB2 crystallites by rocking ± 1 . After 3678 s at 1268 ± 3 C, the sample was cooled to room temperature in <10 min, and XRD patterns confirmed that ZrB2 phase had been completely oxidized and m-ZrO2 was the only crystalline phase present. Similar isothermal studies were also conducted for the ZrB2-SiC composite samples, although at different temperatures and for different durations of times. Each HTXRD pattern was recorded using a 30 s exposure of the sample to the X-ray beam. In the case of 70:30 ZrB2 -SiC sample fifty-two HTXRD patterns were collected with a time resolution of 1 min 16 s at                       1276 ± 3 C, while forty-seven HTXRD patterns were collected with a time resolution of 1 min 27 s for the 50:50 ZrB2-SiC sample at 1334 ± 2 C. After collecting HTXRD patterns for 3847 s and 3776 s for the 70:30 ZrB2-SiC and the 50:50 ZrB2-SiC samples, respectively, the samples were cooled to room temperature in <10 min, and XRD patterns were recorded to ascertain the final crystalline phase compositions after oxidation. Using the Rietveld method, the final compositions of the oxidized 70:30 ZrB2-SiC sample expressed in weight % was 42.7 ± 0.7% ZrB2 , 12.4 ± 0.5% SiC (ICDD PDF # 01-074-2307), 41.6 ± 0.9% mand 3.3 ± 0.2% t-ZrO2 . The oxidized 50:50 ZrB2 -SiC ZrO2 sample was comprised of 34.1 ± 0.4% ZrB2 , 24.9 ± 0.5% SiC, 36.8 ± 0.6% m-ZrO2 and 4.1 ± 0.1% t-ZrO2 by weight. The tZrO2 phase found in both the composite samples after cooling was in contrast to the pure ZrB2 sample. It is believed that most of the t-ZrO2 phase formed at higher temperatures would have achieved significant crystallite coarsening and would transform to m-ZrO2 upon cooling. However, both the composite samples also had ZrB2 phase remaining at the end of the isothermal study, which could have oxidized during cooling to produce small crystallites of t-ZrO2 which remained stable at room temperature. This is consistent with the appearance of t-ZrO2 phase as the first oxidized crystalline phase observed during the non-isothermal studies on ZrB2 reported in Section 3.3. In addition, a boron carbide phase (B13C2 , ICDD PDF # 04-002-9582) was also present as a minor phase in both the cooled composite samples, but this phase was excluded from the final quantitative analysis. The boron carbide phase was not observed in any of the XRD patterns which were acquired at high temperatures. Therefore, it is believed that it may have been formed as a result of rapid cooling of a microstructure which comprised of B2O3 liquid phase in the presence of the SiC phase. Further investigation of this observation was not pursued as it was considered beyond the scope of the current work. The post-oxidation microstructures of the ZrB2 and the ZrB2-SiC composites after the isothermal studies is presented in Fig. 8. The pure ZrB2 sample was completely oxidized and the resultant microstructure was porous throughout (Fig. 8a). The  \\x0c', 'P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  2383  Fig. 8. SEM micrographs of the cross-section of the isothermally oxidized samples: (a) ZrB2 , (b) 70:30 ZrB2 -SiC, and (c) 50:50 ZrB2 -SiC. Please note: the pure diboride sample (ZrB2 ) was oxidized at 1268 ± 3 C for 3678 s; the 70:30 ZrB2 -SiC sample was oxidized at 1276 ± 3 was oxidized at 1334 ± 2 C for 3847 s; and the 50:50 ZrB2 -SiC sample C for 3776 s. All the samples were cooled to room temperature after the completion of the isothermal studies in <10 min.           non-oxidized ZrB2 sample was already porous, and conversion of ZrB2 to m-ZrO2 phase upon oxidation is associated with a further increase of 14.74% in molar volume. The presence of elongated grains of m-ZrO2 phase was a characteristic feature of the final microstructure suggesting preferential grain growth along certain directions in the ZrO2 phase. The microstructure of the oxidized 70:30 ZrB2-SiC and the 50:50 ZrB2 -SiC samples (Fig. 8(b) and (c)) was multilayered. A thin and glassy, Si-rich outside surface layer covered a porous inner layer which was formed above the underlying ZrB2-SiC composite. There was no indication of elongated grain growth in the porous layer of the composite samples, and it was mainly comprised of the ZrO2 phase, as verified by elemental analysis using EDS. Both the oxidized 70:30 ZrB2-SiC and the 50:50 ZrB2 -SiC samples showed the presence of cracks in the porous layer which could be delamination due to thermal expansion mismatch between  the surface layer, the porous oxidized layer, and the underlying ZrB2-SiC matrix. No cracking was observed in the pure ZrB2 sample after oxidation.  3.5. Effect of SiC addition on oxidation of ZrB2 in ZrB2 -SiC composites  The effect of SiC addition on the oxidation of ZrB2 in ZrB2-SiC composites can be gauged quantitatively by following the crystalline phase evolution during high temperature isothermal studies. ZrB2 , t-ZrO2 and Pt, were the only crystalline phases present in the pure ZrB2 sample, while SiC was the additional crystalline phase observed in the ZrB2-SiC composite samples throughout the isothermal study at high temperatures. Results from Rietveld analysis of selected isothermal HTXRD patterns of the samples are presented in Fig. 9. The normalized  Fig. 9. Crystalline phase changes with time during oxidation of (a) ZrB2 in air at 1268 ± 3 air at 1334 ± 2 C.        C, (b) 70:30 ZrB2 -SiC in air at 1276 ± 3     C, and (c) 50:50 ZrB2 -SiC in  \\x0c', '2384  P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  molar concentrations of ZrB2 and t-ZrO2 are presented on the yaxis on the left, while the SiC concentration, normalized by the SiC content in the non-oxidized sample, is plotted on a second y-axis on the right for the ZrB2 -SiC composite samples. Using the QLF the samples were heated to the desired temperature in 5 min and the HTXRD pattern acquisition using the CIP detector was immediately started. The first HTXRD pattern was acquired with a delay as small as 35 s (for the 70:30 ZrB2 -SiC sample) and no larger than 198 s (for the pure ZrB2 sample) from the time the sample(s) had reached the desired temperature(s). Despite of the high porosity in the samples, a parabolic region in the oxidation of ZrB2 was observed for all the samples over the first 500 s, and subsequently, a linear oxidation trend was observed. The slope of the ZrB2 (or t-ZrO2 ) curves in the linear region represents the rate of oxidation of the ZrB2 phase. The SiC content in the 70:30 ZrB2-SiC and the 50:50 ZrB2-SiC samples decreased slowly, and approximately 25% of the SiC phase had oxidized in both the samples in 1 h. The larger error bars for SiC content in the 70:30 ZrB2-SiC sample are due to higher noise in the HTXRD datasets collected for this sample. Based on this analysis it can be qualitatively seen that the oxidation rate of the ZrB2 decreased with increasing SiC content in the samples. Since the samples had different initial porosities, and the isothermal runs for each of the samples reported in this study were conducted at different temperatures, a rigorous comparison of the rate of oxidation of ZrB2 was not possible. It should however be noted that a higher slope of the t-ZrO2 content, suggesting a faster rate of oxidation, was observed in the linear region between 500 and 3000 s for the 70:30 ZrB2-SiC than for the 50:50 ZrB2-SiC. Not only did the 70:30 ZrB2 -SiC have the least porosity and the 50:50 ZrB2 -SiC have the highest porosity of all the samples investigated as part of this study, the isothermal study for 50:50 ZrB2-SiC was also conducted at a markedly higher temperature (1334 ± 2 C). These observations raise some relevant questions regarding the factors that can influence oxidation kinetics in the experimented temperature range. The oxidation protection offered by the silica layer in the examined temperature range in a rigid but porous sample, along with the role of microstructure and porosity in facilitating the formation of the silica layer, certainly deserve some consideration. Perhaps, it could be interesting to explore oxidation behavior of a ZrB2-SiC composite which has a thin layer of SiC on the outside surface. The fractional conversion of ZrB2 to ZrO2 in the studied UHTC diborides, as determined using HTXRD, can also be used to estimate the thickness of the oxidation layer as it is formed. The HTXRD experiments in transmission mode examined a volume element of the sample which extended through the sample and had two surfaces exposed to air with an area ‘A’, determined by the X-ray beam cross-section (see Fig. 2). The fraction of the ZrB2 phase which is converted to t-ZrO2 , therefore, is equivalent to the ratio of the volume of the oxide layer (formed on both exposed surfaces) to the total sampled volume. The final expression for oxide layer thickness, denoted by ‘\\x01t’ and presented in Eq. (2), takes into consideration the fraction of the sample which is porous ‘p’ at the onset of the experiment, the initial volume fraction of ZrB2 (VZrB2 ), and the fractional increase in volume     Fig. 10. Change in thickness of the oxidation layer formed with time in pure  ZrB2 and ZrB2 -SiC composite samples during isothermal studies at tures ≥1250 C.     tempera due to conversion of ZrB2 to t-ZrO2 (\\x01vZrB2→t -ZrO2 ): × (1 − p) × (1 + (x × \\x01vZrB2→t -ZrO2 )) 2 × A  \\x01t = x×t×A×VZrB2  (2)  In the above expression ‘x’ is the molar fraction of ZrB2 converted to t-ZrO2 , which is determined using the HTXRD, and ‘t’ is the thickness of the non-oxidized sample. Any changes in porosity in the sample during oxidation or the contribution of the thin glassy silica surface layer to the oxidation layer thickness are also ignored. An inherent assumption in the derivation of this expression is the movement of the oxidation front into the specimen, parallel to the surface of the specimen. While this may be expected for dense, bulk diboride samples, it can be modified by the presence of porosity, and in particular interconnected porosity in the sample microstructure. Despite of these simplifications, the derived expression can provide a reasonable estimate of the average thickness of the oxidized layer being formed during oxidation especially in dense diboride samples. The effect of SiC on the oxidation layer thickness, formed due to oxidation of ZrB2 to t-ZrO2 , is shown in Fig. 10. The results from pure ZrB2 are also included for comparison. Notwithstanding the differences in the physical microstructures of the starting samples, and the differences in temperatures that were used for the isothermal HTXRD runs for each of the samples, it can be seen that increasing the SiC content decreased the rate of oxidation of ZrB2 . The oxide layer thickness determined using this approach should be further adjusted for the increase in volume associated with t-ZrO2 to m-ZrO2 conversion on cooling to room temperature (\\x01v  4.84%), in order to compare the calculated results with the layer thickness determined by ex situ SEM imaging. The thickness of the oxidation layer formed and observed using the SEM (shown in Fig. 8), correlates well with the calculated values (shown in Fig. 10). It is important to acknowledge that this comparison was undertaken merely as an exercise and not as a rigorous proof of the  \\x0c', 'P. Sarin et al. / Journal of the European Ceramic Society 30 (2010) 2375-2386  2385  concept. However, it still presents a novel alternative to estimate the oxidation layer thickness as it is formed during oxidation of ZrB2 and ZrB2 -SiC ceramics, especially for dense samples, as a function of time and temperature with reasonable accuracy.  4. Summary  High temperature X-ray diffraction was successfully used to examine the oxidation of ZrB2 in situ, at high temperatures in air, in pure ZrB2 and ZrB2 -SiC composite ceramics. Increasing the SiC content in the ZrB2 -SiC composites retarded the oxidation of ZrB2 . Both non-isothermal and isothermal HTXRD studies on the UHTC diboride samples were possible due to the improved X-ray detector system and the QLF, which allowed rapid heating of samples in air up to high temperatures (up to 2000 C). HTXRD methods provide useful information, complementary to the conventional TGA/DSC methods, when used to study the oxidation of diboride ceramics. It was possible to identify and quantify any intermediate crystalline phases that were formed during oxidation of ZrB2 , in real time. The presence of concurrent phases, amorphous or crystalline, or simultaneous reactions, were not limiting and the oxidation of ZrB2 phase could be followed independently. In addition, a novel approach to estimate the thickness of oxidation layer formed during oxidation of ZrB2 and ZrB2-SiC composites, in situ at high temperatures, has been proposed. It is based on fractional conversion of ZrB2 to ZrO2 , as determined using HTXRD, and perhaps will be most appropriate for application to oxidation of dense ZrB2 and ZrB2-SiC composites. Development of UHTC materials can benefit immensely from HTXRD studies. It is anticipated that through such investigations, improved insight will be gained into the effect of sample microstructure and composition on ZrB2 oxidation. This can guide the development of improved models to predict the lifetime and performance of UHTCs under operational conditions.     Acknowledgments  P. Sarin and R. P. Haggerty were supported by the AFOSR FA9550-06-1-0386; P.E. Driemeyer and J.L. Bell were supported by the AFOSR FA9550-06-1-0221; and Z.D. Apostolov was supported by the SMART Scholarship by the Department of Defense. The CIP detector was designed and built under an AFOSR DURIP award number FA9550-04-1-0345. The Advanced Photon Source, Argonne National Laboratory is supported by the US DOE, BES-Materials Sciences, under Contract No: W-31-109-ENG-38.  References  1. Opeka MM, Talmy IG, Zaykoski  JA. Oxidation-based materials  selec tion  for  2000     C + hypersonic  aerosurfaces:  theoretical  considerations 2004;39(19):  and  historical  experience.  Journal  of Materials  Science  5887-904.  2. Monteverde F, Bellosi A, Guicciardi S. Processing and properties of zirco nium diboride-based composites. 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},{
  "_id": 103,
  "PDF": "In situ study of the oxidation of ZrB2 and ZrB2-SiC materials by monitoring the LIF signal of BO2 radicals.pdf",
  "Text": "['Corrosion Science 148 (2019) 31-38  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  In situ study of the oxidation of ZrB2 and ZrB2-SiC materials by monitoring the LIF signal of BO2 radicals  T  V. Guérineaua, A. Julian-Jankowiaka,⁎, G. Vilmartb, N. Dorvalb  a DMAS, ONERA, Université Paris-Saclay, F-92322 Châtillon, France b DPHY, ONERA, Université Paris-Saclay, F-91123 Palaiseau, France  A R T I C L E  I N F O  A B S T R A C T  Keywords: A. Ceramic C. High temperature corrosion C. Oxidation B. Spectroscopy B.Laser induced fluorescence  1.  Introduction  In this study, the Laser Induced Fluorescence (LIF) technique is used to detect BO2 radicals in the gas phase above heated ZrB2 and ZrB2-20 vol. % SiC samples in air, in order to provide an in situ and real-time monitoring of their thermal oxidation. Samples are heated up to 1923 K in air flow with a 2 kW CO2 laser. The BO2 fluorescence and the laser transmission signals are monitored throughout the temperature ramp. This technique allows to detect the key steps of the oxidation (silica formation,...) and to propose more precise oxidation mechanisms as a function of temperature.  ZrB2 and HfB2-based ceramics belong to the Ultra-High Temperature Ceramics (UHTC) family and are good candidates for applications in extreme environments such as thermal protection systems on hypersonic aerospace vehicles or atmospheric re-entry vehicles. The interest brought by these materials is due to their good thermal, mechanical and chemical properties: high melting point, high hardness, chemical inertness and relatively good oxidation resistance in severe environments. The oxidation resistance of ZrB2, HfB2, ZrB2-SiC and HfB2-SiC mastudied [1-10], terials has been thoroughly and the oxidation mein air are relatively well known up to 1873-2073 K. Bulk chanisms ZrB2, when oxidised, develops on the surface a boria (B2O3) layer that protects the material up to 1373 K [11]. Beyond this temperature, B2O3 volatilisation is significant, and the boria layer does not protect the material anymore. The addition of SiC has proven to enhance the oxidation resistance properties up to 1873 K. Indeed, the passive oxidation of SiC (onset at 1473-1573 K [12]) generates SiO2(l), which is able to form, in the presence of boria, a borosilicate glassy layer on the surface. The borosilicate layer is more viscous and more stable than a boria layer, and is therefore able to decrease the diffusion of oxygen at higher temperatures. The glassy layer formed on the surface of UHTC during oxidation is known to be a key point in the oxidation resistance of UHTC. Several attempts have been made to increase the viscosity of the borosilicate glass by increasing liquid immiscibility [13,14]. In  addition, several studies deal with the oxidation behaviour at temperatures higher than 2273 K of either monolithic materials or UHTC matrix composites [15-17] and have shown the formation of a dense zirconia or hafnia layer. However, in most cases, properties of the glassy layer are inferred from post-test analyses and despite the crucial role of the borosilicate glass, few studies attempted to investigate the in situ, and thus continuous, behaviour during a complete oxidation test [14,15]. The thermogravimetric analysis (TGA) [5] is a common method to study the oxidation of UHTC, but mass variations as a function of time only give insights to the sum of the mass gains and mass losses. Sevastyanov et al. [18] have measured the spontaneous emission signals from excited B and Si atoms above an HfB2-SiC sample oxidised at very high temperatures, but limited to the 1873 K and above temperature range. Optical emission spectroscopy was also used by Playez et al. [19] to detect the boron species (B, BO and BO2) in the gas phase at the surface of ZrB2-SiC samples in high-temperature plasma oxidation conditions. Though, these techniques are of great interest, the elements have to be in an excited state and in sufficient quantity to be detected, thus making these techniques slightly sensitive. Nevertheless, reducing the number of samples and experiments and better understanding what happens during oxidation are key points for materials development and optimisation. In this study, the authors have used a laser-based spectroscopic method, the laser induced fluorescence (LIF), to detect gaseous BO2(g) radicals above the sample during oxidation generated by the  ⁎ Corresponding author. E-mail addresses: vincent1guerineau@gmail.com (V. Guérineau), aurelie.jankowiak@onera.fr (A. Julian-Jankowiak), gautier.vilmart@centrale-marseille.fr (G. Vilmart), nelly.dorval@onera.fr (N. Dorval).  https://doi.org/10.1016/j.corsci.2018.11.032 Received 13 July 2018; Received in revised form 11 October 2018; Accepted 26 November 2018  Available online 04 December 2018 0010-938X/ © 2018 Elsevier Ltd. All rights reserved.  \\x0c', 'V. Guérineau et al.  volatilisation of the boria and borosilicate layers. As excitation of species is performed using a dedicated laser with a specific wavelength adapted to the selected BO2(g) specie, the LIF technique allows to detect BO2 radicals even at low concentrations thanks to its higher sensitivity and with a better spatial resolution than emission spectroscopy. The LIF signal of BO2 is monitored through the well-known A2Πu-X2Πg electronic transition with laser excitation tuned to resonance with the (0,0,0-0,0,0) band near 547 nm [20]. To the authors’ knowledge, this paper deals with the first application of the LIF technique to investigation of the oxidation of UHTC samples via the BO2 radicals in the gas phase. Thus, this paper is dedicated to the adaptation of the LIF technique in order to study the real time oxidation behaviour of the glassy layer formed on the surface of ZrB2 and ZrB2-SiC UHTC materials during oxidation in dry air, from 1173 to 1873 K. Finally, the results obtained with LIF measurements are then compared to TGA results.  2. Materials and apparatus  2.1. Materials  To study the interest of LIF diagnostic during the oxidation test, two compositions of UHTC materials were selected: ZrB2 and ZrB2 + 20 vol. % SiC, labelled ZS. The ZrB2 samples were prepared by cold-pressing of as-received ZrB2 powder (H.C. Starck, grade A, d50 = 2.8 μm) at 148 MPa for 10 s, and then sintered at 1823 K for 1 h in an argon atmosphere in a graphite furnace (Pinette Emidecau). Resulting ZrB2 pellets (13 mm in diameter and 5 mm in thickness) have an average porosity of 46%. These samples were prepared as reference without SiC for the LIF experiments. Thus, the sintering of these samples was not studied. ZrB2-SiC samples were prepared from powders of ZrB2 and SiC (H.C. Starck, BF12, d50 = 0.6 μm). Powders were attrition-milled in cyclohexane using zirconia media, dried and sieved. Then, the mixtures were sintered by Spark Plasma Sintering (SPS) in the MATEIS laboratory (FCT, HD 125) at 2323 K for 2 min. Detailed description of the manufacturing of the SiC-containing samples is reported elsewhere [15]. Finally, pellets (15 mm in diameter and 2 mm in thickness) are used for characterisation tests. The bulk density and open porosity of materials were measured by the Archimedes’ method. Then, the densification level was calculated as the ratio of the apparent density to the theoretical density of the powder mixture. Samples characteristics are summarised in Table 1.  2.2. LIF experiment  The LIF method is used to detect gaseous BO2 in the evaporation plume of a sample heated by a 2 kW CO2 laser in air. BO2 is excited by a tunable dye laser from its ground state X 2Πg (0,0,0) to the excited state A 2Πu (0,0,0) at 547.3 nm (in vacuum) (Fig. 1). The excited radical relaxes to its ground electronic state by emitting photons at different wavelengths according to the vibration transitions A 2Πu (0,0,0) X 2Πg (v1,v2,v3) which constitute the dispersed fluorescence spectrum. The transition couple (laser excitation and fluorescence detection) chosen in the LIF experiments is shown in Fig. 1. The fluorescence spectrum measured, after laser excitation at 547.3 nm, is presented in Fig. 1. This spectrum is measured above the surface of a heated ZrB2 sample in an argon/dry air (77/23 vol. %) mixture at 53.3 kPa. It is composed of the  Corrosion Science 148 (2019) 31-38  vibrational bands between 450-660 nm of the electronic system A 2Πu X 2Πg of BO2 [20]. The (0,0,0) - (1,0,0) red-shifted band centred at 580 nm is chosen for the fluorescence light detection. It should be mentioned that the resonant band at the same wavelength as the laser one is not recorded here in order to avoid saturation of the detector with laser scattering from the boron oxide particles formed in the evaporation plume. The second harmonic of a Nd:YAG laser (Quantel, YG781) at 532 nm is used to pump a tunable narrow-linewidth dye laser (Quantel, TDL50), which operates with Rhodamine 575 diluted in methanol, in order to generate a laser beam around 547 nm. Dye laser wavelength is measured by means of a wavelength-meter (WS6-600, Highfiness). The the A 2Πu laser wavelength is finely tuned (1 pm steps) to the peak of (0,0,0) - X 2Πg (0,0,0) band of BO2 at 547.312 nm in vacuum (Fig. 1). The pulse duration is 6 ns, with a repetition rate of 10 Hz, the output energy is 15 mJ/pulse and the spectral linewidth is 0.07 cm−1. The excitation laser beam passes through a diaphragm (3 mm in width) and is focused with a spherical lens (f = 600 mm) to obtain a spot size of 1 mm at the chamber centre. The energy is reduced to 1 mJ/pulse using neutral density filters. The laser energy fluctuations are controlled by means of a photodiode (DET210, Thorlabs). During a thermal oxidation run, the laser energy is recorded at the chamber exit by a photodetector (PD10, Ophir) in order to monitor the beam attenuation through the evaporation plume (absorption pathway in Fig. 2). Fluorescence at 580 nm from BO2, after laser excitation at 547 nm, is collected at right angle from the probe laser beam axis (Fig. 2) by a 100 mm focal-length lens coupled to an optical fibre (1 mm in diameter). Fluorescence is filtered by a monochromator (Jobin-Yvon H20 visible, 200 mm, f/4.2, 4 nm/mm) equipped with a 1 mm wide exit slit, resulting in a 4 nm wide bandpass. The signal is then detected by a visible photomultiplier tube (XP2017B, Photonis) with a 2.5 ns rise time. The output signal from the photomultiplier is time-integrated over 34 ns and averaged over 10 laser shots by a Boxcar system (SR250, Stanford Research Systems). The laser system, oscilloscope and Boxcar gate are triggered at 10 Hz using a delay generator (DG535, Stanford Research Systems). A Labview® interface is used to acquire integrated fluorescence signals via an acquisition PCI card (1.25 MHz, 16 bit) at 10 Hz. The fluorescence band detection (580 nm) is far enough from the laser excitation line to avoid interference by laser scattering. It is assessed by carefully checking the drop in signal when the laser wavelength is tuned out of the excited band during an oxidation test. The probe volume is estimated to be roughly 0.5 mm3. A chamber has been adapted for LIF experiments. This chamber is equipped with 3 horizontal optical windows with a broadband visible antireflection coating (400-700 nm). Two of them are lined up to allow the laser beam propagation through the chamber (Fig. 2). The third one, perpendicular to the other two, is used for the detection of the fluorescence signal. On the top of the chamber, a ZnSe lens is used to focus the CO2 laser beam. The sample is positioned on an alumina support itself placed on a vertical translation stage. Thus, the height of the sample can be adjusted under the focused laser beam which is fixed. The surface temperature is measured through a bichromatic pyrometer working in the 1273-2773 K range (modline 6R-2565, Ircon). The chamber is pumped down to 5 Pa, measured with a Pirani gauge. Inside the chamber, the gases are then injected through a ring having several holes surrounding the sample. The flow rates of gases are regulated by mass flow regulators (Full scale: 10 L.min−1). The gases  Table 1  Sintering conditions, open porosity and oxidation conditions of the studied UHTC materials.  Composition  Sintering process  Applied pressure  Thermal treatment  Open porosity  Densification rate ρ/ρth  ZrB2 ZrB2 + 20 vol. % SiC  Pressureless sintering SPS  -  28 MPa  1823 K, 1 h 2323 K, 2 min  46 % 0.1 %  54 % > 99 %  32  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  Fig. 1. Selected excitation/detection scheme of BO2 (left) and dispersed fluorescence spectrum measured above a heated ZrB2 sample in an Ar/dry air atmosphere (right). Pressure: 53.3 kPa; temperature: between 1623 K and 1773 K during measurement; resolution: 4 nm.  without any optical interference, the wavelength of the laser excitation or of the detection can be modified and the temperature is stabilised during these manipulations. Differentiation between the LIF signal and possible interference signals is described in the Appendix A. Finally, in order to perform Scanning Electron Microscope (SEM) examinations after the oxidation tests, samples are impregnated in an in half, polished down to 3 μm and a carbon layer is epoxy resin, cut then deposited on the surface. Identification of the nature and thicknesses of each layer of the oxidised layers was performed using an SEM (DSM 982, Zeiss) equipped with Energy Dispersive Spectroscopy (EDS). For comparison, some TGA experiments were performed in a SETSYS Evolution 16/18 (SETARAM) to measure mass variations during oxidation. Parallelepipeds, 2 × 2 x 4 mm3 in size, were extracted from sintered samples. Heating rate was 2 K/min, and tests were performed from room temperature to 1823 K with a 50 ml/min air flow.  3. Results  3.1. TGA measurement  tests  The SiC-containing sample has been submitted to TGA, and the results are presented in Fig. 3. The ZS mass gain begins at 1113 K. This is attributed to the oxidation of ZrB2 into ZrO2 and liquid B2O3(l) (reaction 1). This is in quite a good agreement with Sarin et al. [21] and Monteverde et al. [5] studies who have observed a mass gain starting around 973 K for a ZrB2 sample in air using the same kind of TGA measurements. This slope starts decreasing at around 1373 K. This is due to the fact that at this temperature, the volatilisation of B2O3 is less and less compensated by the formation of ZrO2 and B2O3(l), and therefore the mass gain starts increasing more slowly (reaction 2). Then, the slope increases at 1473 K, due to the onset of the passive oxidation of SiC. This will generate SiO2(l) that will form, with the boria glassy layer, a borosilicate layer, thus preventing the volatilisation of B2O3(l) and inducing a mass gain (less mass loss). Then, the SiO2-B2O3 glassy layer protects the sample from the oxidising atmosphere and mass gain decreases (less mass gain) (reactions 3,4). This explains why the slope of the mass gain increases first (formation of the borosilicate layer) and then decreases (hindering oxidation through a more stable and viscous glassy layer). However, at 1673 K, the mass gain increases noticeably, and this is due to an increased volatilisation of B2O3, which enhances the oxidation.  Fig. 2. Laser Induced Fluorescence (LIF) setup.  used are synthetic air (Alphagaz 1) and argon (Argon N48 ARCAL PRIME). Absolute pressure is measured with a Baratron gauge. This pressure is regulated around a set value of 0.1 MPa by a throttle valve (235B, MKS) and a regulator (651, MKS). This system allows working in dynamic conditions: gases are injected in the chamber, and evacuated at a rate which depends on the closing rate of the throttle valve.  2.3. Experimental protocol  The sample is positioned on the holder at the centre of the chamber. The alignment of the excitation laser beam is checked in order to target just above the centre of the sample. Vacuum is made in the chamber, and then the gases are injected inside the chamber. Every sample was oxidised under 0.1 MPa of synthetic air. Once the chamber pressure attains 0.1 MPa, the CO2 laser is turned on and heating of the sample surface begins. Heating of the sample is achieved by increasing manually the power of the CO2 laser. The reason for a manual control of the laser power is that, several times during an experiment, temperature has to be stabilised (the temperature ramp is temporarily stopped) in order to make several adjustments. For instance, when the fluorescence of BO2 is detected for the first time during an experiment, the height of the sample is adjusted in order to optimise the signal intensity. Moreover, at higher temperatures, to ensure that the detected signal originates from BO2,  33  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  Fig. 3. Mass variation vs. temperature for ZS sample (2 K/min temperature ramp). The TGA is performed under a 50 ml/min air flow.  ZrB2 + 5/2 O2 → ZrO2(s) +B2O3 (l)  B2O3(l) → B2O3(g)  SiC + 2O2 → SiO2(l) + CO2(g)  SiO2 (l)+B2O3 (l) → {SiO2:B2O3}(l)  (1)  (2)  (3)  (4)  TGA measurements provide some details of the oxidation phenomenon. However, mass variation analysis of such a binary system is not so easy and, mass loss and gain due to the volatilisation of B2O3 and the formation of SiO2, respectively, cannot be precisely determined. Moreover, the composition of the glassy phase cannot be assessed using this technique. The LIF experiments detailed in the following sections, show that the LIF is a promising technique to obtain more information species,…) about oxidation phenomena (reactions, volatile and the evolution of the composition of the glassy layer as a function of temperature.  3.2. LIF signal of BO2 radicals during oxidation of ZrB2 samples  ZrB2 samples were oxidized under 0.1 MPa dry air and two kinds of tests were performed: one test where the LIF signal is measured and another one for which the laser light scattering (Rayleigh scattering or Mie scattering) generated by the very fine boron oxide particles is measured (Fig. 4). The laser scattering signal is measured by tuning the laser excitation at 560 nm (away from the bands of BO2) and the detection window is shifted up by 3 nm from the laser wavelength in order to monitor the tail of the laser scattering band. Thereby, it was possible to discriminate between the LIF and laser scattering signals. In both tests, the laser transmission is simultaneously monitored. Fig. 4a presents the evolution of the surface temperature, the intensity of the LIF signal and the transmitted signal of the excitation laser as a function of time for ZrB2 sample in air, under 0.1 MPa. The laser power starts to increase at t = 0, but the pyrometer only measures temperature from 1263 K. Three temperatures of interest can be noticed: at 1333 K, the LIF signal appears. It slowly increases and then, at 1523 K, the LIF signal increases brutally up to 1573 K, where the LIF signal drastically drops in spite of the continuous temperature increase. The laser transmission starts dropping at 1373 K, to finally reach a very low level in the end (0.1). The appearance of the LIF signal at 1333 K suggests that BO2(g) comes from the boria layer, as ZrB2 samples, like ZS samples, tend to develop a stable external boria layer at those temperatures (Fig. 4) [21]. However, BO2(g) may come from oxidation of glassy or gaseous boria. Thermodynamically, oxidation of B2O3(g) is favoured compared to the oxidation of B2O3(l). Therefore, B2O3(g) is first formed by volatilisation from the boria glassy layer (reaction 2), and then, it may either  Fig. 4. (a) Surface temperature (left axis), LIF and transmission signals (right axis) vs. time and (b) Surface temperature (left axis), Rayleigh (or Mie) scattering and transmission signals (right axis) vs. time of a ZrB2 sample during an oxidation test in dry air under 0.1 MPa. Inset in (b): zoom on LIF and Rayleigh (or Mie) signals between 1273 and 1573 K.  decompose or oxidise (reactions 5,6, Gibbs energies are calculated with HSC Chemistry v8.0 software (Outotec)):  B O (g)  2  3  2B O (g)  2  3  BO (g)  BO (g) ;  2  ΔG  +  →  154.300 *  T  561273  J  +  = −  O (g)  2  →  4 BO (g) ;  2  ΔG  +  169.800 *  T  541510  J  +  = −  (5)  (6)  Whether it is reaction 5 or reaction 6 that is privileged, what really matters is that the formation of BO2 is correlated to the volatilisation of the boria layer. The decrease in transmission means that the atmosphere above the sample is more opaque, and is caused by the evaporation of B2O3(l) into B2O3(g) which condensates into droplets, thus impeding the transmission of the excitation laser. Moreover, Tripp et al. [11] have shown that the oxidation behaviour of ZrB2 would change  34  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  between 1373 and 1473 K, from parabolic to linear behaviour, due to the volatilisation of the vitreous phase at the surface. This is in agreement with the transmission decrease observed at 1463 K, which is an indirect proof of the massive volatilisation of B2O3(g). However, it is difficult to differentiate whether the LIF signal peak is only due to an increase in the fluorescence signal of BO2(g), or an increase in the laser induced incandescence (LII) signal from the B2O3(l) droplets that can interfere with the LIF signal (Appendix). Once the transmission is too low, the measured signal decreases because the laser energy is too much attenuated. This behaviour is in agreement with the catastrophic volatilisation of B2O3, leading to the disappearance of the vitreous phase at the surface. In order to differentiate the part of signal coming from BO2(g) and B2O3(l) droplets, another test has been performed in the same conditions where the Rayleigh (or Mie) scattering (560 / 563 nm) is recorded. As can be noticed in Fig. 4b, the laser scattering signal presents some differences as a function of time compared to the LIF signal. First, the curve is much less smooth, and many oscillation peaks of the signal are observed, similarly to the laser transmission curve. These variations result from the flow oscillations caused by the exhaust throttle valve regulation (pressure regulation system), which are then detected through the measurements of the laser transmission and scattering signals. These signals are directly dependent on the concentration of particles released in the plume above the sample. The decrease in the transmission signal begins around 1473 K. Non-surprisingly, the drop in transmission starts at the same temperature as the appearance of Rayleigh scattering signal, around 1483 K, whereas the LIF signal was first detected at 1333 K. As Rayleigh scattering originates from the nascent condensed B2O3(l) particles, this signal is thus only detected once the evaporation of B2O3(l) is strongly activated. Retrospectively, it confirms that the signal observed at 1333 K, and its progressive increase, is a pure LIF signal of gaseous BO2, without any interference. Moreover, it can be noticed that the LIF signal is very smooth at low temperature and exhibits noticeable fluctuations around 1523 K, thus emphasising the condensation into B2O3 droplets. Based on these considerations, the authors assume that the signal detected during the oxidation experiments was mostly due to the lightinduced fluorescence of BO2. Therefore, the LIF signal variations are interpreted as a consequence of the volatilisation of the boria or borosilicate layers.  3.3.  Influence of SiC on the LIF signal  The analysis of the oxidation behaviour of ZS will allow studying the influence of SiC. The LIF signal as a function of time for a ZS sample is shown in Fig. 5. A general trend is that the experiments with SiCcontaining materials are much longer than with pure ZrB2. This is due to the fact that a higher CO2 laser power must be applied to reach high  Fig. 5. Surface temperature (left axis), LIF signal and laser transmission (right axis) vs. time of a ZrB2-SiC sample during an oxidation test in dry air under 0.1 MPa. Key steps numbered to are described in the text.  35  to  temperatures. This can be explained by the fact that SiC-containing samples are very dense, large, and the SiC presence allows the formation of a more stable glassy phase at the surface which is less isolating than the zirconia layer underneath. Four key steps are spotted in Fig. 5, and are numbered from . The LIF signal is first observed at 1263 K ( ). Between and the LIF signal slowly increases. The step at 1563 K, is characterised by an inversion of the trend: the LIF signal decreases, despite the continuous increase in temperature. The LIF signal keeps decreasing up to step (1643 K) where the LIF signal enters into what could be called the \"dynamic phase\". This phase, which lasts until step is characterised by small signal bumps, immediately followed by a decrease in the signal. Finally, a high intensity LIF signal peak and a concomitant drastic drop in the laser transmission are observed in step (1863 K). This step is very similar to what was observed for ZrB2, but it happens at a higher temperature. The test is stopped at this time because the excitation laser beam is too much absorbed, and the LIF signal is no longer detected. Thus, the presence of SiC has a strong influence on the LIF signal as a function of time. In order to clarify the analysis of this signal, several key steps (labelled from a) to e)) are depicted in Fig. 6. When the LIF signal appears ( , a)), it means that there is a B2O3(l) vitreous phase at the surface, which allows the formation of a detectable quantity of BO2(g), likely through reactions 2 and 5 or 6. At this temperature (1263 K), SiC is not yet oxidised (Fig. 3). This is confirmed by SEM examinations in several studies in which SiO2 is not detected below 1473 K [4]. The following steady increase in the LIF signal intensity is due to the increased volatilisation of B2O3, leading to an increased quantity of BO2(g). Step is essential and emphasises a clear difference compared to the tests performed on ZrB2 samples. This decrease in the LIF signal intensity happens at a temperature where SiC starts to oxidise passively [4,22-24] through reaction 3. Therefore, the formation of SiO2(l) will feed the vitreous phase, thus forming a borosilicate layer which is more viscous and more stable than a boria layer [1,5,6,12,32,33]. As the layer is more stable, less LIF signal is observed, due to the volatilisation of its constituents, despite the continuous temperature increase. Actually, the volatilisation of B2O3(l) may be hampered by the decrease in its activity in the B2O3SiO2 system [25]. Moreover, Karlsdottir et al. [26-28] have reported that when a borosilicate layer is subjected to high temperatures, a preferential evaporation of B2O3 is observed. This will form a concentration gradient (visible at (b)), and the B2O3 concentration will be, in proportion, lower on the surface than at the bottom of the glassy layer. At the very top of the surface, this lower concentration of B2O3 also explains the LIF signal decrease. In order to explain the onset of steps and , a technical bias caused by the experimental setup must be explained. When increasing the CO2 laser power is required, the set value is modified and sometimes, some power peaks are observed, superior to the ordered power setpoint. Those power peaks trigger sudden temperature increases (as an example, a 40 K rise in 20 s is observed). Step , characterised by a sharp increase in the LIF signal, is triggered by a power peak, as well as the signal bumps between steps and . Before step (and before the power peak), the borosilicate layer is in equilibrium between volatilisation and refilling of its constituents. Kinetics studies [22,24,29] show that at these temperatures, the oxidation kinetics of ZS is parabolic, which means that the glassy layer is protective towards oxidation. The power peak breaks this equilibrium and leads to a quick volatilisation of the top of the glassy layer (c)). This reveals the B2O3-rich glassy layer, where the activity of B2O3 is higher and therefore the volatilisation of B2O3 is favoured. The glassy layer is therefore less thick, less stable, and offers less protection to the material underneath: a BO2(g) LIF signal peak is observed in consequence. The glassy layer being destabilised, the material underneath is oxidised more intensely, and therefore quickly reforms a stable glassy layer (d)), and the LIF signal decreases. As the temperature is  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  Fig. 6. Evolution of the boria and borosilicate layers interpreted from the LIF signal vs. time. Steps labelled from a) to e) are described in the text.  continuously increasing, the average intensity of the signal also increases. The subsequent signal bumps have the same origin (power peak) and the same consequences (second c), d) steps). The very high intensity signal peak, that characterises the onset of step , is also triggered by a power peak. But at this high temperature, the volatilisation of the glassy layer is very intense, and the active oxidation of SiC is favoured via reaction 7:  SiC + O2(g) = SiO(g) + CO(g)  (7)  Consequently, the material can no longer compensate the loss of the glassy layer, and a catastrophic volatilisation is therefore observed (e)). This catastrophic volatilisation is characterised by the LIF signal peak and the drop in the laser transmission signal, which means that there is an intense formation of B2O3(l) and SiO particles in the atmosphere. After the test, SEM/EDS analyses allow to visualise the state of the glassy layer at the surface (Fig. 7). The surface of the ZS sample after oxidation at 1873 K shows the presence of a residual glassy layer, but also the presence of many recrystallised zirconia particles, grains and dendrites. According to Karlsdottir and Halloran [26,30], the presence  of ZrO2 structures in the borosilicate is explained by the following mechanism: during the oxidation of ZS, a BSZ (boria-silica-zirconia) liquid phase is formed. The solubility of zirconia is greater in B2O3 than in SiO2. When this BSZ liquid emerges at the surface, the preferential evaporation of B2O3 will lower its concentration in the BSZ liquid and will cause an increase in the viscosity of the glass and crystallisation of ZrO2. The BSZ liquid will move into the two-phase region with solid ZrO2 and liquid BSZ. This is in agreement with the LIF observations. The observation of a LIF signal throughout the test means that B2O3 is constantly evaporating from the glassy layer. Moreover, the test ended with a catastrophic volatilisation of the glassy layer, which likely removed a great portion of the remaining B2O3. Those conditions favour the saturation of the solubilised zirconia, which therefore recrystallises. The surface and transverse section micrographs of oxidised ZS (Fig. 7d, Fig. 8b) confirm the observation of a catastrophic volatilisation of the glassy layer: the zirconia layer is bare at some places, while at others a thin glassy layer is observed. A SiCdepleted layer is also observed, i.e. a layer where SiC, only, was actively oxidised, so that only empty pores are left in the ZrB2 matrix. Because of  Fig. 7. SEM micrographs of different areas of the surface of a ZS sample oxidised at 1873 K. In grey, the borosilicate layer and, dendritic structures. Evidence of the recrystallisation of ZrO2 (a, b, c) and of complete evaporation of the glassy layer (d).  in white, zirconia inclusions and  36  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  Fig. 8. SEM micrographs of transverse sections of an oxidised ZS sample (1873 K). The different layers are: (1) a borosilicate glass, (2) the ZrO2 layer, (3) the SiCdepleted ZrB2 layer and (4) the bulk material. In (b) the glassy layer has totally disappeared through evaporation.  the oxygen gradient in the oxidised layer, SiO(g) formed by reaction 7 can react with oxygen via reaction 8 and therefore refill the glassy layer [25,31].  SiO(g) + ½ O2(g) = SiO2(l)  (8)  The beneficial influence of SiC on the oxidation resistance of UHTC is well known [1,5,6,12,32,33]. The detection of the LIF signal of BO2(g) allows to demonstrate, during the oxidation process, that, when SiO2(l) was added to the glassy layer, this latter was more stable at higher temperatures, prevented the volatilisation of B2O3 and was able to maintain its integrity even after several power peaks.  3.4. Comparison with TGA results  The LIF signal is first detected at 1263 K, whereas the mass gain observed in TGA starts at 1113 K: this means that LIF detection is less sensitive than TGA for detecting the beginning of oxidation. This is due to the fact that the LIF signal appears when there is enough BO2(g) to be detected, therefore after the oxidation has started. Then, at 1563 K, the LIF signal decreases despite the continuous temperature increase. This has been attributed to the influence of the oxidation of SiC. In the TGA in the slope at 1573 K, which is curve, there is a slight decrease therefore interpreted as the influence of SiC, which would overall hinder the oxidation of the sample. If these two methods are compared, then it can be noticed that the LIF detection is much more sensitive to a change in the chemistry of the glassy layer on the surface, in the present case: the introduction of SiO2 in the boria layer in order to form a more protective borosilicate layer. At 1643 K, it was shown that the LIF signal entered into the “dynamic phase”. The temperature itself is not very important, since it depends on the time, or on the temperature, at which the power peak happens. In the TGA study, at 1673 K, there was a sharp mass gain, attributed to an enhanced volatilisation of B2O3 which in turn, enhance the oxidation. The “dynamic phase” observed would, through LIF measurements does illustrate this state. At those temperatures, there are both an intense volatilisation of the surface glass and an intense oxidation underneath, which continuously replenish the glassy layer. The glassy layer on the surface is therefore metastable: if the temperature ramp is slow enough, then the glassy layer is rather stable on the surface. However, if there is a sudden increase in the temperature, the volatilisation is enhanced (LIF signal increases) because the glassy layer is destabilised. Nevertheless, as the oxidation is still intense, and the temperature peak quickly ceases, the glassy layer is quickly refilled (LIF signal decreases). The LIF detection allows to embrace more accurately the oxidation features on the surface than the TGA, because this latter only measures the sum of the mass gains and losses.  4. Conclusion  ZrB2 and ZrB2-SiC materials were oxidised up to 1873 K in a dry air atmosphere, and their oxidation behaviours were studied through the measurement of BO2(g), via the LIF method, above the sample during  oxidation. Thanks to this method, key steps of the oxidation can be detected for both ZrB2 and ZrB2-SiC materials: (i) the starting point of the influence of oxidation of SiC in the stabilisation of the surface glassy layer and (ii) the catastrophic volatilisation of the glassy layer. The LIF method also allowed to emphasise the “dynamic phase”, between 1623 and 1873 K, which is characterised by the formation of a metastable glassy borosilicate layer on the surface, in equilibrium with an intense volatilisation of its components and its refilling by an intense oxidation of the material underneath. The dynamic behaviour of this glassy phase was evidenced through the power peaks, which allowed detecting the temporary destabilisation of this layer. Therefore, LIF is a promising method, which can be quantitative, to study in situ, and thus continuously, the oxidation behaviour of UHTC, and could be a good alternative to the classical post-test and ex situ studies, which are very consuming in terms of time and samples.  Data availability  The raw/processed data required to reproduce these findings cannot be shared at this time due to legal or ethical reasons.  Authors’ contributions  V. Guérineau is the PhD student who has manufactured the samples, tested them, performed the post-test analyses and written the article. N. Dorval and G. Vilmart are the specialists of the LIF technique and have helped to modify the bench and prepare the experiments. V. Guérineau, N. Dorval and G. Vilmart have performed all the LIF experiments. A. Julian-Jankowiak is the PhD supervisor, she has helped to modify the bench, to analyse the results and to write the paper. All the authors have read and approved the final version of article.  this  Declaration of interest  None.  Acknowledgements  The authors would like to thank M. Bejet and T. Schmid (ONERA) for modifying and improving the oxidation chamber, Dr P. Beauchêne (ONERA) for his kind and efficient assistance and for very fruitful discussions.  Appendix A  The signals, possibly interfering with the 580 nm after a laser excitation at 547 nm, are:  LIF  signal  of BO2  at  The black body radiation of the heated surface and of the hot gases and hot particles in the evaporation plume produces an emission continuum. Moreover, the spontaneous emission from the excited  37  \\x0c', 'V. Guérineau et al.  Corrosion Science 148 (2019) 31-38  [16]  [13]  [14]  [21]  [20]  experience, J. Mater. Sci. 39 (2018) 5887-5904, https://doi.org/10.1023/B:JMSC. 0000041686.21788.77 (n.d.). E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2and HfB2-based ultra-high temperature ceramics: effect of Ta additions, J. Mater. Sci. 39 (2018) 5969-5977, https://doi.org/10.1023/B:JMSC.0000041693.32531.d1 (n.d.). I.G. Talmy, J.A. Zaykoski, M.M. Opeka, High-temperature chemistry and oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3, and TaSi2, J. Am. Ceram. Soc. 91 (2008) 2250-2257, https://doi.org/10.1111/j.1551-2916.2008.02420.x . [15] V. Guérineau, A. Julian-Jankowiak, Oxidation mechanisms under water vapour conditions of ZrB2-SiC and HfB2-SiC based materials up to 2400°C, J. Eur. Ceram. Soc. 38 (2018) 421-432, https://doi.org/10.1016/j.jeurceramsoc.2017.09.015 . C. Carney, A. Paul, S. Venugopal, T. Parthasarathy, J. Binner, A. Katz, P. Brown, Qualitative analysis of hafnium diboride based ultra high temperature ceramics under oxyacetylene torch testing at temperatures above 2100 °C, J. Eur. Ceram. Soc. 34 (2014) 1045-1051, https://doi.org/10.1016/j.jeurceramsoc.2013.11.018 . [17] H. Li, X. Yao, Y. Zhang, K. Li, L. Guo, L. Liu, Effect of heat flux on ablation behaviour and mechanism of C/C-ZrB2-SiC composite under oxyacetylene torch flame, Corros. Sci. 74 (2013) 265-270, https://doi.org/10.1016/j.corsci.2013.04.052 . [18] V.G. Sevastyanov, E.P. Simonenko, A.N. Gordeev, N.P. Simonenko, A.F. Kolesnikov, E.K. Papynov, O.O. Shichalin, V.A. Avramenko, N.T. Kuznetsov, Behavior of a sample of the ceramic material HfB2-SiC (45 vol %) in the flow of dissociated air and the analysis of the emission spectrum of the boundary layer above its surface, Russ. J. Inorg. Chem. 60 (2015) 1360-1373, https://doi.org/10.1134/ S0036023615110133. [19] M. Playez, D.G. Fletcher, J. Marschall, W.G. Fahrenholtz, G.E. Hilmas, S. Zhu, Optical Emission spectroscopy during plasmatron testing of ZrB2-SiC ultrahightemperature ceramic composites, J. Thermophys. Heat Transf. 23 (2009) 279-285, https://doi.org/10.2514/1.39974. J.W.C. Johns, The absorption spectrum of BO2, Can. J. Phys. 39 (1961) 1738-1768, https://doi.org/10.1139/p61-199. P. Sarin, P.E. Driemeyer, R.P. Haggerty, D.-K. Kim, J.L. Bell, Z.D. Apostolov, W.M. Kriven, In situ studies of oxidation of ZrB2 and ZrB2-SiC composites at high temperatures, J. Eur. Ceram. Soc. 30 (2010) 2375-2386, https://doi.org/10.1016/j. jeurceramsoc.2010.03.009. [22] W.C. Tripp, H.H. Davis, H.C. Graham, Effect of a SiC addition on the oxidation of ZrB2, ResearchGate. 52 (1973). https://www.researchgate.net/publication/ 255363900_Effect_of_a_SiC_addition_on_the_oxidation_of_ZrB2 (Accessed October 13, 2016). [23] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500 °C, J. Eur. Ceram. Soc. 27 (2007) 2495-2501, https://doi.org/10.1016/j.jeurceramsoc.2006.10.012 . [24] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364, https://doi.org/ 10.1111/j.1551-2916.2007.01583.x . E.P. Simonenko, D.V. Sevast’yanov, N.P. Simonenko, V.G. Sevast’yanov, N.T. Kuznetsov, Promising ultra-high-temperature ceramic materials for aerospace applications, Russ. J. Inorg. Chem. 58 (2013) 1669-1693, https://doi.org/10.1134/ S0036023613140039. S.N. Karlsdottir, J.W. Halloran, Oxidation of ZrB2-SiC: influence of SiC content on solid and liquid oxide phase formation, J. Am. Ceram. Soc. 92 (2009) 481-486, https://doi.org/10.1111/j.1551-2916.2008.02874.x . S.N. Karlsdottir, J.W. Halloran, C.E. Henderson, Convection patterns in liquid oxide films on ZrB 2-SiC composites oxidized at a high temperature, J. Am. Ceram. Soc. 90 (2007) 2863-2867, https://doi.org/10.1111/j.1551-2916.2007.01784.x . S.N. Karlsdottir, J.W. Halloran, Formation of oxide scales on zirconium diboride-silicon carbide composites during oxidation: relation of subscale recession to liquid oxide flow, J. Am. Ceram. Soc. 91 (2008) 3652-3658, https://doi.org/10. 1111/j.1551-2916.2008.02639.x . S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra-high temperature ceramics for aeropropulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767, https://doi.org/10.1016/S0955-2219(02)00140-1 . S.N. Karlsdottir, J.W. Halloran, A.N. Grundy, Zirconia transport by liquid convection during oxidation of zirconium diboride-silicon carbide, J. Am. Ceram. Soc. 91 (2008) 272-277, https://doi.org/10.1111/j.1551-2916.2007.02142.x . T.A. Parthasarathy, R.A. Rapp, M. Opeka, M.K. Cinibulk, Modeling oxidation kinetics of SiC-containing refractory diborides, J. Am. Ceram. Soc. 95 (2012) 338-349, https://doi.org/10.1111/j.1551-2916.2011.04927.x . J. Han, P. Hu, X. Zhang, S. Meng, Oxidation behavior of zirconium diboride-silicon carbide at 1800 °C, Scr. Mater. 57 (2007) 825-828, https://doi.org/10.1016/j. scriptamat.2007.07.009. [33] D.D. Jayaseelan, E. Zapata-Solvas, C.M. Carney, A. Katz, P. Brown, W.E. Lee, Microstructural evolution of HfB2 based ceramics during oxidation at 1600-2000°C, Adv. Appl. Ceram. 114 (2015) 277-295, https://doi.org/10.1179/1743676115Y. 0000000018.  [25]  [26]  [27]  [28]  [29]  [30]  [31]  [32]  BO2 (chemiluminescence) may interfere with LIF of BO2. In the absence of the laser beam excitation source, no signal was detected, thus demonstrating that these interferences are negligible under these experimental conditions. The laser light scattering from the fine boron oxide particles which density increases in the plume with the heating temperature. In this case, the elastic scattering (Mie or Rayleigh scattering) is very intense close to the laser wavelength but vanishes for sufficiently shifted fluorescence bands which is the case for the selected redshifted fluorescence band of BO2 at 580 nm compared to the 547 nm excitation line. Another possible interference is the Laser-Induced Incandescence (LII) signal produced by the incandescence of nano-sized particles (< 100 nm) after absorption of the laser beam at 547 nm. This signal generates a broad continuous spectrum without any particular structure such as a black body radiation. This signal can become intense at high temperatures and can impede the detection of BO2(g) fluorescence when the density of nanoparticles in the smoke becomes sufficiently high. Indeed, a persistent signal is still detected when the detection window is moved out any fluorescence bands. This phenomenon may be attributed to the laser incandescence background since a white smoke is observed above the sample.  We achieved in accurately detecting the BO2 fluorescence with a good signal/noise ratio up to 1773-1873 K. At higher temperatures, the laser beam energy is strongly attenuated by the rise of the density of boron oxide smoke, and this prevents a proper measurement of LIF signal. The temperature limit depends on the sample.  References  [5]  [3]  [1] C.M. Carney, P. Mogilvesky, T.A. Parthasarathy, Oxidation behavior of zirconium diboride silicon carbide produced by the spark plasma sintering method, J. Am. Ceram. Soc. 92 (2009) 2046-2052, https://doi.org/10.1111/j.1551-2916.2009. 03134.x. [2] W.G. Fahrenholtz, G.E. Hilmas, Oxidation of ultra-high temperature transition metal diboride ceramics, Int. Mater. Rev. 57 (2012) 61-72, https://doi.org/10. 1179/1743280411Y.0000000012 . P.A. Williams, R. Sakidja, J.H. Perepezko, P. Ritt, Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content, J. Eur. Ceram. Soc. 32 (2012) 3875-3883, https://doi.org/10.1016/j.jeurceramsoc.2012.05.021 . [4] X.-H. Zhang, P. Hu, J.-C. Han, Structure evolution of ZrB2-SiC during the oxidation in air, J. Mater. Res. 23 (2008) 1961-1972, https://doi.org/10.1557/JMR.2008. 0251. F. Monteverde, The thermal stability in air of hot-pressed diboride matrix composites for uses at ultra-high temperatures, Corros. Sci. 47 (2005) 2020-2033, https:// doi.org/10.1016/j.corsci.2004.09.019 . [6] C.M. Carney, Oxidation resistance of hafnium diboride—silicon carbide from 1400 to 2000 °C, J. Mater. Sci. 44 (2009) 5673-5681, https://doi.org/10.1007/s10853009-3799-7. [7] K. Shugart, E. Opila, SiC Depletion in ZrB2-30 vol% SiC at Ultrahigh Temperatures, J. Am. Ceram. Soc. 98 (2015) 1673-1683, https://doi.org/10.1111/jace.13519. F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC composite, J. Eur. Ceram. Soc. 25 (2005) 1025-1031, https://doi.org/10.1016/j.jeurceramsoc. 2004.05.009. [9] K. Shugart, W. Jennings, E. Opila, Initial stages of ZrB2-30 vol% SiC oxidation at 1500°C, J. Am. Ceram. Soc. 97 (2014) 1645-1651, https://doi.org/10.1111/jace. 12843. P. Hu, W. Guolin, Z. Wang, Oxidation mechanism and resistance of ZrB2-SiC composites, Corros. Sci. 51 (2009) 2724-2732, https://doi.org/10.1016/j.corsci. 2009.07.005. [11] W.C. Tripp, H.C. Graham, Thermogravimetric study of the oxidation of ZrB2 in the temperature range of 800° to 1500°C, J. Electrochem. Soc. 118 (1971) 1195-1199, https://doi.org/10.1149/1.2408279 . [12] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000°C + hypersonic aerosurfaces: theoretical considerations and historical  [10]  [8]  38  \\x0c']"
},{
  "_id": 104,
  "PDF": "In Situ Synthesis of Ultrafine ZrB2–SiC Composite Powders and the Pressureless Sintering Behaviors.pdf",
  "Text": "['In Situ Synthesis of Ultrafine ZrB2-SiC Composite Powders and the  Pressureless Sintering Behaviors  Yongjie Yan,  z,y  Hui Zhang,  z,y  Zhengren Huang,  w,z  Jianxue Liu,  z  and Dongliang Jiang  z  z  Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China  y  Graduate School of the Chinese Academy of Sciences, Beijing 100039, China  Ultrafine ZrB2-SiC composite powders have been synthesized in situ using carbothermal reduction reactions via the sol-gel method at 15001C for 1 h. The powders synthesized had a relatively (o200 smaller average crystallite size nm), a larger specific surface area (B20 m2/g), and a lower oxygen content (B1.0 wt %). Composites of ZrB2120 wt% SiC were pressureless sintered to B96.6% theoretical density at 22501C for 2 h under an argon atmosphere using B4C and Mo as sintering aids. Vickers hardness and flexural strength of the sintered ceramic composites were 13.970.3 GPa and 294714 MPa, respectively. The microstructure of the composites revealed that elongated SiC grain dispersed uniformly in the ZrB2 matrix. Oxidation from 11001 to 16001C for 30 min showed no decrease in strength below 14001C but considerable decrease in strength with a rapid weight increment was observed above 15001C. The formation of a protective borosilicate glassy coating appeared at 14001C and was gradually destroyed in the form of bubble at higher temperatures.  I.  Introduction  ULTRAHIGH-TEMPERATURE ceramics, with a combination of excellent properties, such as a high melting point, high hardness, high thermal and electrical conductivity, and good  oxidation resistance and chemical attack resistance, have been  proposed as potential candidates  for applications such as fur nace elements, plasma-arc electrodes, rocket engines, and thermal protection structures for space vehicles.1-4  ZrB2-SiC ceramic composites, of known to have better strength and oxidation resistance monolithic ZrB2 materials.2,5 The enhances the oxidation resistance of ZrB2 by the formation of a protective borosilicate glass layer.6 High strength has been  suitable composition, are  than  addition  of  SiC particles  attributed to maintaining a fine grain size and a uniform distriphase.5 Monteverde  bution  of  the  reinforcing  and  Bellosi  reported that incorporation of ultrafine SiC improved the sinterability and mechanical properties of ZrB2.7,8 Rezaie et al. proposed that the strength-limiting flaws in ZrB2-SiC were the larger SiC particles in the microstructure.9 Subsequently, Zhu et al.10 also investigated the influence of SiC particle size on the  microstructure and mechanical properties of ZrB2-SiC ceramics. It was shown that starting SiC particle of smaller sizes led to  improved densification, finer grain sizes, and higher  strength.  Thus, a fine grain size and a uniform distribution of SiC parti cles played an important  role in the sintering and mechanical  properties of the ZrB2-based ceramics.  Generally, densification of ZrB2-SiC ceramics typically requires temperatures of 19001C or higher and applied pressures of 20 MPa or more.11,12 Sintering aids including MoSi2, Ni, TiB2, B4C, Si3N4, Al2O3, and Y2O3 have been introduced to improve the sinterability of ZrB2-based ceramics.13-15 However, fabrication of complex components requires expensive and time consuming diamond machining. Development of  the pressure less sintering process would be advantageous as components can  be  fabricated to a near-net  shape.  In this  sense, pressureless  sintering may be an effective way to extend the applications of  the ZrB2-based ceramics. In our former work, ultrafine ZrB2 powders have been synthesized successfully via the sol-gel method, and the pressureless sintering behaviors of the ZrB2-based ceramics were studied.16,17 It is known that ultrafine or nanosized SiC powders have been synthesized using this method for many years.18, 19 The mecha nisms can be described according to the following two equations:  ZrO2 þ B2O3 þ 5C ! ZrB2 þ 5CO \"  (1)  SiO2 þ 3C ! SiC þ 2CO \"  (2)  In  addition, we  found  that  the  two  reactions  occurred  thermodynamically at 1723 and 1768 K calculated by the free  energy changes using the dynamical data from NIST-JANAF tables.20 Hence,  it provides a possibility to synthesize  in situ  ultrafine ZrB2-SiC composite powders. ultrafine ZrB2-SiC composite powders were in situ synthesized using inorganic precursors ZrOCl2 \\x01 8H2O, H3BO3, and silica sol as sources of silica, boron oxide, and silicon, respectively. The  In the present  study,  pressureless sintering process, mechanical properties, and oxida tion resistance of the ceramic composites were mainly studied.  II.  Experimental Procedure  (1)  Powder Synthesis and Characterization  Zirconium oxychloride (Sinopharm Chemical Reagent, Shanghai,  China) and boric acid (Shanghai Experiment Reagent, Shanghai,  China), used as the sources of zirconia and boron oxide, were of  analytical grade. Silica sol  (JN-30, Heyuan Chemical, Guang dong, China) was used as the source of silica and the content of silica was B27 wt%. Phenolic resin (PF-04872, Henan Bond  Chemical, Luoyang, China) was used as whose char yield was B50 wt%.  the  carbon source,  The amounts of zirconium oxychloride, boric acid, silica sol,  and phenolic resin were determined according to reactions (1)  and (2) by the composition of ZrB2-20 wt% SiC. It is mentioned that H3BO3 was 10 wt% excessive according to our former work. ZrOCl2 \\x01 8H2O and H3BO3 in ethanol. Silica sol was added and Firstly, an inorganic hybrid solution was prepared by dissolving the pH value was adjusted to 4 using dilute ammonia. Then,  phenolic resin was added to form a mixed sol and dilute ammonia  was  sequentially added to promote the gelation process. The  N. Padture—contributing editor  This work was financially supported by the National Natural Science Foundation of  China.  w  Author  to  whom  correspondence  should  be  addressed.  e-mail:  zhrhuang@  mail.sic.ac.cn  Manuscript No. 23739. Received September 14, 2007; approved December 4, 2007.  Journal  J. Am. Ceram. Soc., 91 [4] 1372 - 1376 (2008)  DOI: 10.1111/j.1551-2916.2008.02296.x  r 2008 The American Ceramic Society  1372  \\x0c', 'treatment of the gel and the heat-treatment process were similar to the synthesis of ZrB2 reported in our former work.16 The phase compositions were determined using X-ray  diffractometry (XRD, RAX-10A, Rigaku, Tokyo, Japan) with CuKa radiation. The different element contents were determined  by chemical analysis. The crystallite size and morphology were  characterized by a field emission scanning electron microscope  (JSM-6700F, JEOL, Tokyo, Japan), and the surface areas were  characterized  by  nitrogen  adsorption-desorption  isotherms  obtained  using  an  automatic  adsorption  instrument  (ASAP  2010, Micromeritic, Atlanta, GA) at 77 K.  (2)  Pressureless Sintering Process  Molybdenum powders (MP-1, Luoyang Highteck, Luoyang, China) had a purity of 99.95% with particle size o50 mm, and B4C (F1200, Jingangzuan Boron Carbide, Mudanjiang, China) had a purity of 90% with particle size o3 mm. Certain amounts of in situ synthesized ZrB2-20 wt% SiC powders with 4 wt% Mo, 2 wt% B4C, and 2wt% phenolic resin were ball milled using SiC milling media in ethanol for 24 h. After milling, ethanol was  removed using rotary evaporation (Rotavapor R-124, Bucchi, Flawil, Germany) at 701C. The composite powders were firstly at B60 MPa  uniaxially  dry  pressed  using a steel die and 4 mm \\x02 6 mm \\x02 followed by cold isostatic 40 mm at B200 MPa for 120 s. The green compacts were heat treated to 8001C (2 h in vacuum) to remove the binder at a rate of 51C/min. After binder removal, the samples were presintered to 16001C (2 h in vacuum) and then sintered to 22501C (2 h in argon) at a rate of 101C/min. After sintering, the furnace  pressing  into  was cooled to room temperature under flowing argon. All  the  heat-treatment processes were carried out in a high-temperature  graphite  resistance  furnace  (High-Multi  10,000, Fujidempa  Kogyo, Osaka, Japan).  (3) Mechanical, Oxidation Resistance, and Characterization  Bulk density was measured using the Archimedes displacement  method and theoretical density was based on the rule of mixture.  Vickers hardness and three-point flexural  strength were mea sured by Vickers indentation (Model 300, Tukon, Canton, MA)  using a load of 5 kg and a dwell  time of 10 s and a universal  tester (Instron-1195, Instron, Canton, MA) over a 30 mm span  with a cross-head speed of 0.5 mm/min. The microstructure of  the sintered body was observed by electron probe microanalysis  (EPMA) (JXA-8100, JEOL) with energy-dispersive spectroscopy. Each side of the bars with dimensions of 3 mm \\x02 4 mm \\x02 36 mm was polished to a 1 mm finish using diamond abrasives for  furnace oxidation. Before oxidation, specimens were cleaned in  an ultrasonic bath in acetone, ethanol, and deionized water and then dried at 1201C for 3 h. The bars were placed in Al2O3 crucibles and heat treatment was carried out in a box-type resis tance  furnace (HT 40/17, Nabertherm GmbH, Lilienthal, from 11001 to 16001C for 30 min at every 1001C  Germany)  interval. After oxidation,  the flexural strength and weight gain  were measured. For microstructural analysis of the oxide scale mm finish  by EPMA,  cross  sections were  polished  to  a  1  perpendicular to the top surface of the oxidized bars.  III.  Results and Discussion  XRD patterns of the precursor powders at different heattreatment temperatures from 11001 to 16001C for 1 h are shown in Fig. 1. At 11001C, m-ZrO2, t-ZrO2, and B2O3 phases are observed and t-ZrO2 is the dominant phase. The carbon present in amorphous. At 12001C, the sample remains the intensity of  m-ZrO2 and t-ZrO2 decreases and a ZrB2 phase has already appeared. An SiC phase appeared at 14001C and the crystal form was identified as b-SiC according to the location and in tensity of the peaks (PDF#291129). However, we can see that the ZrC phase is still present at 15001C, which is believed to be  the result of the volatilization of B2O3 at high temperatures because of its low melting point. XRD patterns at 15001 and  16001C show no significant  change, which indicates  that  the  carbothermal reduction process was basically complete. The rel ative weight  fraction of SiC was calculated to be about 18.6%  (slightly lower than the theoretical composition) by the relative  intensity ratio method using a Guinier-Ha¨ gg camera (Expectron Instrument, Solna, Sweden) with CuKa1 indicated by \\x0315001C radiation and Si as an internal standard, the ZrC phase still appeared in the XRD  XDC-1000,  Jungner  in Fig.  l. However,  patterns and more precise analysis should be conducted in fur ther studies.  Figure 2 shows the SEM image of the in situ synthesized ZrB2-SiC composite powders at 15001C for 1 h. Particle sizes mainly distribute over 100-200 nm with a spherical morphology  and several agglomerations appear. However,  it  is difficult  to  distinguish the difference between ZrB2 and SiC. Some properties of the composite powders, such as the par ticle size, BET surface area, and different element contents (wt %),  are listed in Table I. It can be seen that is B20m2/g, which is much larger  the BET surface area  than some  commercially  available ZrB2 or SiC powders. ratio calculated is 0.54, higher than 0.5, and the Si/C molar ratio  In addition,  the Zr/B molar  is 0.95,  lower than 0.1, which indicates the existence of ZrC in  accordance with the XRD results. Considering the free carbon  and the oxygen contents, we chose B4C as the suitable sintering aid in order to remove the residue oxygen and promote the  carbothermal reduction process according to Eq. (3):  4ZrO2 þ B4C þ 3C\\x00!4ZrB2 þ 4CO \"  (3)  Fig. 1. X-ray diffractometry patterns of the precursor powders different heat-treatment temperatures from 11001 to 16001C for 1 h.  at  Fig. 2.  Scanning electron microscopic image of the in situ synthesized  ZrB2-SiC composite powders.  April 2008  Communications of the American Ceramic Society  1373  \\x0c', 'The mechanical properties of the ZrB2-SiC ceramic composites sintered at 22501C for 2 h are listed in Table II. The relative  density of  the composite is considerably high, and it  indicates  that  the in situ synthesized composite powder has good sinter ability. The bending strength and Vickers hardness of the composite is 294714 MPa and 13.970.3 GPa, respectively, which  are lower compared with the values measured for the hotpressed composites.11,12 Fracture toughness and elastic modulus were calculated using the following formulas21:  K IC ¼ 0:026  \\x12  E 1=2P1=2a  C3=2  \\x13  (4)  where P is the indentation load, a the half-length of the indent,  C the half-length of the crack, and E the elastic modulus of the  composite. E is  calculated from the  elastic modulus of  the  components, according to the  rule of mixtures. The  fracture  toughness of the ZrB2-SiC ceramic composite was about 5.34.6870.14 MPa \\x01 m1/2, which was higher than or comparable to the hot-pressed material.22  E ¼ 3LðP2 \\x00 P1 Þ 2bh2 ðe2 \\x00 e1 Þ \\x02 10\\x003  (5)  where P1, P2 are the loads, e1, e2 span, b the width of the specimen, and h the thickness of  the strains, L the length of  the  specimen. The calculated elastic modulus of the ZrB2-SiC composites is 31579 Gpa, which is lower than the values calculated  according  to  the  rule  of mixtures.  (EZrB2 5 450 GPa  and  ESiC 5 414 GPa).  The microstructure of  the ZrB2-SiC ceramic composite was characterized by examining polished cross sections in Fig. 3. It  was shown that SiC (dark phase) were regularly dispersed in the  ZrB2 (gray phase) matrix and appeared to be distinctly elongated (approximately 1-5 mm wide by 5-20 mm long). The rod-like SiC grains are thought to be caused by the transition from b-SiC to a-SiC, but the mechanisms should be studied further. An SEM  image from a fracture surface of the ZrB2-SiC ceramic composite is shown in Fig. 4. The fracture propagation mode is mainly  intragranular and the composite had low residual porosity.  Figure 5 shows the weight gain and flexural strength with an increase in temperature from 11001 to 16001C every 1001C for  30 min. The weight gain increased slowly from 1.07 to 3.34 mg/ cm2 below 14001C and the flexural strength showed nearly no  change 15001C,  compared with  the  room-temperature  strength. At  the weight gain accelerated and the flexural  strength  decreased to 80% of the room-temperature strength, which in dicated that the material had been destroyed to some extent at temperature approached 16001C,  this  temperature. When the  the flexural strength rapidly decreased to 121 MPa.  Generally, exposure of ZrB2-SiC to air at high temperatures results in stoichiometric oxidation to ZrO2 (cr) and SiO2 (l ) by reactions (6) and (7):  4ZrB2 ðcrÞ þ 5 2  O2 ðgÞ\\x00!ZrO2 ðcrÞ þ B2O3 ðl Þ  (6)  aSiCðcrÞ þ 3 2  O2 ðgÞ\\x00!SiO2 ðcrÞ þ COðgÞ  (7)  Table I. Particle Size, BET Surface Area, and Different Element Contents (wt%) of the ZrB2-20 wt% SiC Composite Powders Synthesized at 15001C for 1 h  Particle  size (nm)  BET  (m2/g)  Zr  (%)  Si  (%)  B  (%)  CTotal  (%)  CFree  (%)  O  (%)  o200  B20  66.54  9.05  14.57  6.33  2.26  1.19  Table II.  Mechanical Properties of the ZrB2-SiC Ceramic Composites Sintered at 22501C for 2 h  Composition  (wt%)  Relative  density  (%)  Bending  strength  (MPa)  Vickers  hardness  (GPa)  Fracture  toughness (MPa \\x01 m1/2)  Elastic  modulus  (GPa)  ZrB2-20 SiC  97.6  294714  13.970.3  4.6870.14  31579  Fig. 3. Back-scattered electron image of the polished surface of ZrB2-SiC ceramic composite at 22501C for 2 h.  the  Fig. 4. Scanning electron microscopic image of the fracture surface of the ZrB2-SiC ceramic composite at 22501C for 2 h.  Fig. 5. Dependence of weight gain and flexural strength of composites on temperature from 11001 to 16001C for 30 min.  1374  Communications of the American Ceramic Society  Vol. 91, No. 4  \\x0c', 'April 2008  Communications of the American Ceramic Society  1375  Fig. 6.  Scanning electron microscopic images of the cross-sectional oxidized ZrB2-SiC materials from 11001 to 16001C for 30 min.  Several authors have reported that oxidation of ZrB2-SiC at 15001C in air produces a structure that consists of four layers:  (1) a continuous silica layer on the surface; (2) a Zr-rich oxidized  layer embedded in amorphous silica; (3) a layer of SiC-depleted ZrB2; and (4) unaffected ZrB2-SiC.23-25  Figure 6 shows the SEM images of the cross-sectional oxidized ZrB2-SiC materials from 11001 to 16001C for 30 min. The red dotted line and blue dotted line separate the different scales the oxide surface. Below 13001C,  the oxidation of SiC was  of  much slower  than that of ZrB2 and the SiC particles did not  \\x0c', '1376  Communications of the American Ceramic Society  Vol. 91, No. 4  oxidize appreciably. The oxide scale mainly consisted of: layer of ZrO2 B10 mm thickness, (2) a layer of SiC-depleted ZrB2 B30 mm thick, and (3) unaffected ZrB2-SiC. In addition, no continuous B2O3 (l) layer was found to form above the ZrO2 layer even at 11001C. At 13001C, the ZrO2 scale enlarged to B15mm and some flakes appeared, which was caused by volume  (1) a  expansion upon conversion of ZrB2 to ZrO2 and B2O3. As the temperature approaches 14001C, the vapor pressure of  B2O3 increases substantially, leading to its rapid evaporation. In addition, SiC starts to oxidize, producing molten SiO2 and gaseous species such as CO in this temperature regime. A thin and smooth SiO2-rich layer (B5 mm) covered the underlying material and could, potentially, provide a barrier to oxygen diffusion.  Other literatures reported that the SiO2-rich layer was expected to contain some B2O3 and ZrO2 by the formation of a single anionic matrix of the melt according to Eq. (8)26:  B2O3 ðl Þ þ ZrO2 ðcrÞ þ SiO2 ðcrÞ ! kZrO2 \\x01 mB2O3 \\x01 nSiO2  (8)  At 15001C, the thickness of the SiO2-rich layer increased to B10 mm and became uneven. This may be due to wetting char acteristics or other local variations such as composition, surface  topology, or surface cracks that might enhance the local oxidathe SiC-depleted layer enlarged to B80 tion rate. Meanwhile, mm and had a porous structure from which the SiC had been  partially or entirely removed. The structure of the specimen heated to 16001C for 30 min was similar to the structure of the specimen exposed to 15001C, except that the reaction layers were thicker (B30 mm) and the SiC-depleted layer enlarged to B120 mm. The SiO2-rich layer became more uneven and relatively extensive bubble formation was observed on the surface  exposed to the air. The material may be destroyed considerably,  combined with the result of flexural strength degradation.  IV.  Summary  Ultrafine ZrB2-SiC composite powders have been in situ synthesized using carbothermal reduction reactions via the sol-gel  method. The synthesized powders had a larger specific surface  area that can be pressureless sintered to 96.6% theoretical density at 22501C for 2 h due to its favorable sinterability. The sin tered ceramic composites had moderate mechanical properties and maintained their flexural strength to 14001C for 30 min in  air. The protective borosilicate glassy coating appeared at 14001C and was gradually destroyed in the form of bubbles at  higher temperatures.  References  1R. A. Cutler, ‘‘Engineering Properties of Borides’’; pp. 787-803 in Ceramics and  Glasses, Engineered Materials Handbook, Vol. 4, Edited by S. J. Schneider. ASM  International, Materials Park, OH, 1992. 2A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  ‘‘High Strength Zirconium Diboride-Based Ceramics,’’ J. Am. Ceram. Soc., 87 [6]  1170-2 (2004).  3G. V. Samsonov, Refractory Transition Metal Compounds: High Temperature  Cermets. Academic Press, New York, 1964. 4K. Upadhya, J. Yang, and W. Hoffman, ‘‘Materials for Ultrahigh Temperature  Structural Applications,’’ Am. Ceram. Soc. Bull., 76 [12] 51-6 (1997). 5W. C. Tripp, H. H. Davis, and H. C. Graham,  ‘‘Effect of a Silicon Carbide  Addition on the Oxidation of Zirconium Diboride,’’ Am. Ceram. Soc. Bull., 52 [8]  612-6 (1973). 6F. Monteverde and A. 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Guicciardi,  ‘‘Processing and Properties of  Zirconium Diboride-Based Composites,’’  J. Eur. Ceram. Soc.,  22  [3]  279-88  (2002). 12A. Bellosi, F. Monteverde, D. D. Fabriche, and C. Melandri, ‘‘Microstructure  and Properties of ZrB2-Based Ceramics,’’ J. Mater. Process. Manuf. Sci., 9 [10] 156-69 (2000). 13F. Monteverde and A. Bellosi,  ‘‘Effect of the Addition of Silicon Nitride on  Sintering Behaviour and Microstructure of Zirconium Diboride,’’ Scr. Mater., 46  [3] 223-8 (2002). 14R. A. Andrievskii, L. A. Korolev, V. V. Klimenko, A. G. Lanin, I. I. Spivak,  and I. L. Taubin,  ‘‘Effect of Zirconium Carbide and Carbon Additions on Some  Physicomechanical Properties of Zirconium Diboride,’’ Powder Metall. Metal  Ceram., 19 [2] 93-4 (1980). 15S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Pressureless Densificat ion of Zirconium Diboride with Boron Carbide Additions,’’ J. Am. Ceram. Soc.,  89 [5] 1544-50 (2006). 16Y. J. Yan, Z. R. Huang, S. M. Dong, and D. L. Jiang,  ‘‘New Route to Syn thesize Ultra-Fine Zirconium Diboride Powders Using Inorganic-Organic Hybrid  Precursors,’’ J. Am. Ceram. Soc., 89 [11] 3585-88 (2006). 17Y. J. Yan, Z. R. Huang, S. M. Dong, and D. L. Jiang, ‘‘Pressureless Sintering  of High-Density ZrB2-SiC Ceramic Composites,’’ J. Am. Ceram. Soc., 89 [11] 3589-92 (2006). 18V. Raman, O. P. Bahl,  ‘‘Synthesis of Silicon Carbide  and U. Dhawan,  Through the Sol-Gel Process from Different Precursors,’’ J. Mater. Sci., 30 [10]  2686-93 (1995). 19Y. S. Jung, O. J. Kwon, and S. M. Oh, ‘‘Formation of Silica-Coated Carbon Powder and Conversion to Spherical b-Silicon Carbide by Carbothermal Reduc tion,’’ J. Am. Ceram. Soc., 85 [8] 2134-6 (2002). 20M. W. Jr. Chase, NIST-JANAF Thermochemical Tables, 4th edition, Amer ican Institute of Physics, Woodbury, NY, 1998. 21 ‘‘Testing Methods for Fracture Toughness of Fine Ceramics’’; Japanese In dustrial Standard No. JIS R 1607, 1995. 22Y. W. Bao, Z. C. Cao, and Q. R. Ma,  ‘‘Test Methods for Elastic Moduli of  Fine Ceramics  (Advanced Ceramics, Advanced Technical Ceramics)—Bending  Method’’; National Standard of  the People’s Republic of China GB/T 10700,  2006. 23E. J. Opila, S. Levine, and J. Lorincz,  ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39 [19]  5969-77 (2004). 24M. Gasch, D. Ellerby, E.  Irby, S. Bechman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties,  and Arc  Jet Oxidation of Hafnium Diboride/Silicon  Carbide Ultra High Temperature Ceramics,’’  J. Mater. Sci.,  39 [19] 5925-37  (2004). 25W. G. Fahrenholtz,  ‘‘Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 26I. B. Ban’kovskaya and V. A. Zhabrev,  ‘‘Kinetic Analysis of the Heat Resis tance of ZrB2-SiC Composites,’’ J. Glass Phys. Chem., 31 [4] 482-8 (2004).  &  \\x0c']"
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  "_id": 105,
  "PDF": "Influence of sintering temperature, SiC particle size and Y 2 O 3 addition on the densification, microstructure and oxidation resistance of ZrB 2 –SiC ceramics.pdf",
  "Text": "['Journal of the European Ceramic Society 36 (2016) 3041-3049  Contents lists available at www.sciencedirect.com  Journal  of  the  European  Ceramic  Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Influence of sintering temperature, SiC particle size and Y2O3 addition on the densification, microstructure and oxidation resistance of ZrB2-SiC ceramics  Z. Ková ˇcová a , L’ . Ba ˇca a,b,∗ , E. Neubauer c , M. Kitzmantel c  a Department of Ceramics, Glass and Cement, Institute of Inorganic Chemistry, Technology and Materials, Faculty of Chemical and Food Technology, Slovak  University of Technology, Radlinského 9, SK-812 37 Bratislava, Slovakia b Aerospace & Advanced Composites GmbH, Viktor-Kaplan-Strasse 2, A-2700 Wiener Neustadt, Austria c RHP-Technology GmbH, Forschungsund Technologiezentrum, A-2444 Seibersdorf, Austria  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 18 August 2015 Received in revised form 20 December 2015 Accepted 21 December 2015 Available online 2 January 2016  Keywords:  ZrB2 -SiC Y2 O3 Oxidation resistance Microstructure UHTCs  1.   Introduction  The influence of sintering temperature, silicon carbide particle size (50-60 nm, 0.9  \\u242em and 44  \\u242em) and investigated up to 1650   C  yttrium oxide addition on oxidation behaviour of ZrB2 -SiC ceramics was  in static atmosphere. Weight changes were measured and the thickness of formed oxide layer was evaluated by the light and scanning electron microscopy. Results show enhanced oxidation resistance of ZrB2 -SiC composites sintered at 1750   C in spite of their lower density (93.2%) due to the grain growth cessation resulting  in the refinement of microstructure. Similarly the oxidation resistance was also enhanced by refinement of microstructure due the submicron or nano-sized starting SiC powders. The addition of Y2O3 caused desired stabilization of cubic ZrO2 phase; nevertheless  the oxidation resistance of Y2O3 doped samples was  inferior  in comparison to the basic material at oxidation temperatures exceeding 1500   C.  © 2015 Elsevier Ltd. All rights reserved.  Ultra-high temperature ceramics (UHTCs)  is a very  interesting group of materials with melting points over 2000  C. There are more than 300 materials, including the refractory metals (Hf, Zr, Nb, Ir, Ta, W), oxides (HfO2 , ZrO2 , ThO2 ), a variety of transition metal borides, carbides and nitrides, as well as other compounds [1]. They represent a class of promising materials for use in extreme applications: thermal protection systems on hypersonic aerospace vehicles or re-entry vehicles, materials for wing  leading edges and nose tips, as well as specific components  for propulsion,  furnace elements, refractory crucibles, etc., With respect to the application area, high melting temperature  is only one of many other required properties. Mechanical strength, density, thermal conductivity, thermal expansion and also cost are important factors in the material selection process. Nevertheless, oxidation resistance is the major issue for  the application of UHTCs. Several studies have affirmed  that diborides of the group IVb are the most resistant to oxidation [2-4].  ∗ Corresponding author. E-mail address: lubos.baca@stuba.sk (L’ . Ba ˇca).  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.12.028 0955-2219/© 2015 Elsevier Ltd. All rights reserved.  Zirconium diboride (ZrB2 ) is a leading member of UHTCs due to exceptional combination of physical and chemical properties such as high melting point,  low density, high hardness, high  thermal and electrical conductivity, chemical inertness and relatively good oxidation resistance in severe environments [1,4-8]. Compared to HfB2 , advantage of ZrB2 is  its  lower price and theoretical density [9]. It has been shown that addition of SiC to ZrB2 ceramics resulted in significant improvement of oxidation resistance and mechanical properties [3,4,6,8,10,11]. Due to the oxidation starting  from the surface  into the core of composites, a structure with several distinct  layers has been observed [3,4,12-17]. Exposure of ZrB2 -SiC composite to an oxidizing environment at elevated temperatures results in a formation of protective borosilicate glass layer according to following equations [12,13]:  ZrB2 (s) + 5 2  O2 (g) →  ZrO2 (s) + B2O3 (l,g)  SiC (s) +  3  2  O2 (g) →  SiO2 (l) + CO (g)  (1)  (2)  The  oxidation  resistance  of  ZrB2 -SiC  ceramics  can be  also affected  by  the density  of  sintered materials  and/or  by  their microstructure. The influence of sintering parameters on densifica                                             \\x0c', '3042   Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049  Fig. 1. Microstructures of  (1700   C/50 MPa).  samples   sintered   at different   conditions:   (a) ZS2-1   (2100   C/30 MPa),   (b) ZS2-2   (1900   C/30 MPa),   (c) ZS2-3   (1750   C/30 MPa),   (d) ZS2-5  tion and mechanical properties have been studied by Zhao et al. [25] and Zamharir et al. [26]. The applied pressure and the sintering temperature have been  identified as the main  factors controlling the densification of hot-pressed composites. The hardness and fracture toughness of ZrB2 -SiC composites were dependent on the holding time at maximum temperature due to the growth of grains. Furthermore, the densification of composites has been  improved by decreasing the particle size of SiC powders. The reason was that the larger number of small SiC particles was more effective to  inhibit the grain growth of ZrB2 matrix. The grain growth inhibition seems to be one of the main  factors  influencing not only the densification and mechanical properties, but also the oxidation resistance of UHTCs. Moreover,  it has been  reported  that nano-sized SiC particles  improved the oxidation resistance of ZrB2 -SiC ceramics, [21,24] enhanced densification, [18-21] room temperature bending strength and fracture toughness [7,18,20,22,23] although at the  expense of high  temperature strength  [27,28]. The higher  room temperature strength has been attributed to the homogenous distribution of reinforcing phase as well as to  limited grain growth of ZrB2 during densification. However, few studies demonstrated [13,20] that SiC particles dispersed along the grain boundaries of ZrB2 tend to grow during the hot pressing, therefore the sintering parameters have to be carefully selected. Except of grain growth  inhibiting additives also  the sintering additives, most frequently rare earth oxides (REO), have a strong influence on the properties of composites due to the formation of a protective refractory coating on  the surface of UHTCs. Several researches have  investigated the microstructure and mechanical properties of REO-doped ceramics and  reported  that REO have beneficial effect on the densification of UHTCs and their hardness, fracture toughness and creep resistance [29-33].  Fig. 2. Average grain size of ZrB2 and SiC of samples sintered under different conditions.  \\x0c', 'Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049   3043  Fig. 3. Mass change vs. oxidation temperature of the tested samples hot pressed at different sintering temperatures.  The aim of this work was to study the  influence of (i) sintering  temperature,  (ii) SiC particle size and  (iii) Y2O3 addition on densification, microstructure and oxidation resistance of ZrB2 -SiC to 1650  C.  composites up  In detail,  the microstructure  features were carefully analyzed and  the grain size effect after rapid hot pressing on oxidation behaviour was  investigated. The addition of Y2O3 to ZrB2 -SiC ceramics could  lead to the stabilization and crystallization of ZrO2 in dense protective  layer  formed during oxidation at temperatures when borosilicate glass is not more effective.  2. Experimental procedure  Commercially available powders were used for the preparation of 80 wt.% ZrB2 and 20 wt.% SiC composites and their characteristics are listed in Table 1. Three different SiC powders with different particle size (50-60 nm, 0.9  \\u242em and 44  \\u242em) were used to investigate the influence of SiC grain size on grain growth of ZrB2 matrix. 8 wt.% Y2O3 sintering aid was added to the ZrB2 -SiC mixture. The powder mixtures were homogenized on rolls in cyclohexane using zirconia balls for 24 h and dried in evaporator. Discs with 50 mm diameter were  fabricated by rapid hot pressing (DSP518, Dr. Fritsch Sondermaschinen GmbH, Germany) using graphite die with BN coating under various conditions. The sintered samples  Fig. 5. Oxidation surfaces of samples ZS2-1 (sintering T = 2100   C/15 min) and ZS2-3 (sintering T = 1750   C/15 min) after 1 h at 1100   C in static air.  Table 1 Main characteristics of starting powders.  Starting powder   Labelling   Particle size   Supplier  ZrB2 ␤-SiC  ␣-SiC  SiC  Y2 O3  Z  S1  S2  S3  Y   \\u242em   6.1  50-60 nm  0.9  44  <10   \\u242em  \\u242em  \\u242em   ABCR IoLiTec ESK GmbH WesterTonbergau Alfa Aesar  were machined by diamond wheel and then the bulk densities were measured using the Archimedes method in distilled water. All compositions, sintering conditions and relative densities of prepared samples are summarised in Table 2. For example samples ZS2-1 up to ZS2-5 were prepared  from the same starting powder mixture (ZrB2 with 0.9  \\u242em SiC), but under different sintering conditions, Table 2. Small cylinders with diameter 10 mm and height 4-5 mm were cut out from the hot-pressed discs by electrical discharge machining (EDM) and exposed to oxidation. Only samples with densities equal or higher than 90% were used for determination of oxidation resistance. Oxidation tests were carried out in a furnace at temperatures 1100  C, 1300  C, 1500  C and 1650  C for 60 min in stagnant air. Cross-sections of oxidized  samples were embedded  in polymer matrix, grinded and polished with diamond  slurries up  to \\u242em finish. The phase compositions of samples were analysed 1  by X-ray diffraction (Stoe Theta-Theta with CoK␣  radiation). Oxi Fig. 4. The temperature dependence of oxide scale thickness on two different sintering temperatures (1700   C and 1900   C) for samples ZrB2—20 wt.% SiC.  \\x0c', '3044   Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049  Table 2 Summary of sintering conditions and relative densities of prepared samples. Heating rate 100   C/min.  Sample   ZS1  ZS2-1  ZS2-2  ZS2-3  ZS2-4  ZS2-5  ZS3  ZS2Y-1  ZS2Y-2   Sintering conditions   1900   C/15 min/30 MPa  2100   C/15 min/30 MPa  1900   C/15 min/30 MPa  1750   C/15 min/30 MPa  1700   C/15 min/30 MPa  1700   C/15 min/50 MPa  1900   C/15 min/30 MPa 2100   C/15 min/30 MPa  1900   C/15 min/30 MPa   Relative density (%)  97.1 98.7 98.3 93.2 82.4 90.0 96.8 99.6 99.5  dation  resistance was evaluated according  to  the mass changes and  layer  thickness after oxidation. The  specimen  surfaces and microstructure were observed using  light microscope (Leica DMI 5000 M equipped with digital camera DP 25 Olympus and with a soft  imaging system “Scandium” for  image processing and 2D/3D image analysis) and field emission scanning electron microscope (FE SEM, Carl Zeiss SUPRA TM 40VP). The microstructural parameters were determined  from  the SEM micrographs using  ImageJ software with implemented particle analysis.  3. Results and discussion  3.1. Densification  In order to guarantee the shortest possible sintering cycle, the rapid hot pressing technique with a heating rate of 100 K/min was used for the preparation of ZrB2 -SiC composites. The sintering conditions and final densities of samples are summarised  in Table 2. The relative densities of samples increased with sintering temperature. The maximum densities of 98.7% and 98.3% were observed for samples ZS2-1 and ZS2-2 hot-pressed at 2100  C and 1900  C for 15 min, respectively. At lower sintering temperatures (1700  C and 1750  C) densities on the  level of 82% and 93% could be achieved (samples ZS2-3 and ZS2-4). Higher applied pressure during hotpressing (50 MPa) at 1700  C resulted in increase of density by about 8%  (sample ZS2-5). Taking  into account  the economic aspects of sample preparation  (sintering  temperature)  the obtained results show that the optimum sintering temperature for densification of ZrB2 -SiC composites is 1900  C. Spark plasma sintering  (SPS) with the same heating rate was used by Zapata-Solvas et al.  [34]  for  the preparation of 40 mm ZrB2 -SiC disc with SiC particle  size of 0.7  \\u242em. SPS of ZrB2 -SiC ceramic composite at 1950  C for 4 min with a pressure of 50 MPa resulted  in density of 99%.  It  indicates that when submicron SiC powders are used  for  the preparation of ZrB2 -SiC composites a temperature about 1900  C could be sufficient  for densification. However, our study shows that also by using of coarser (44  \\u242em) SiC powder almost  fully dense ZrB2 -SiC composites can be prepared by fast sintering at 1900  C (Table 2). Moreover, the addition of Y2O3 allowed the densification of these UHTCs to 99.5% of theoretical density. This result is consistent with the finding of Zhang et al. [31] where addition of 3 vol.% of Y2O3 improved sinterability of the powders and supressed grain growth by reacting with oxides on the starting powder surfaces. When 2 wt.% La2O3 as an additive was used, the final density of prepared ZrB2 -SiC-La2O3 composite reached 99.6% of theoretical density [35]. From the literature  is known, that  if nanometre sized SiC powders for SPS of ZrB2 -20 wt.% SiC composites were used,  the sintering  temperature can drop down to values as low as 1600 or 1700  C for almost complete densification [18]. These results are consistent with the observation of Zamora et al. [36] where nanosized starting powders were used for densification of ZrB2 by SPS. They showed that,  Fig. 6. Oxidation surfaces of samples ZS2-1 (sintering T = 2100   C/15 min) and ZS2-3 (sintering T = 1750   C/15 min) after 1 h at 1300   C in static air.  the temperature  limit for the SPS densification of ZrB2 decreased continuously with decreasing crystal size  in the starting powder. Except of the small particle size also the higher surface oxide content of nanosized SiC can contribute to the better densification of this composite. This can promote formation of liquid phase during sintering and can assist in densification.  3.2. Oxidation resistance of ZrB2 -SiC composites  3.2.1.   Influence of sintering temperature  Fig. 1a-d  shows  the  scanning  electron micrographs of  the ground and polished surfaces of the ZrB2 -SiC samples sintered at from 1700  C up to 2100  C  temperatures started  for 15 min. The microstructure of the composite is regular with homogenous dispersion of dark SiC particles within  light grey ZrB2 matrix and no agglomeration was detected. The average grain size of SiC and ZrB2 after sintering at respective temperatures was determined and the results summarised in Fig. 2. The average grain size of both constituents  increases with increasing sintering temperature. The difference  is more evident between samples sintered at 1900  C and 2100  C, where in latter \\u242em and case the average grain size of SiC and ZrB2 increased to 2.0  \\u242em, respectively. The difference  3.5  in SiC and ZrB2 grain size  is not so remarkable in the case of samples sintered at lower temperatures (1700  C, 1750  C and 1900  C). At 1900  C only very small increase in grain sizes was observed though the relative density of ZS2-2 sample reached 98.3%.  \\x0c', 'Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049   3045  neous with small particle sizes of SiC as well as ZrB2 . The only explanation of such behaviour can be in the differences in the relative densities, ZS2-5 reached only 90% in comparison with 93.2% for the ZS2-3. Higher porosity can lead to higher degree of boride oxidation and increase of the weight. Contrary, the weight gain obtained for samples ZS2-1, ZS2-2 and ZS2-3 show only small changes and seem to be independent on the relative density. Similar trend was observed also at oxidation temperatures 1300  C and 1500  C. At these temperatures the oxygen transport through the borosilicate glass and mixed ZrO2 -SiO2 layers is responsible for mass gain [13]. The analysis of  the oxidation surface of samples ZS2-3  (finer microstructure, sintered at 1750  C) and ZS2-1 (coarser microstructure, sintered at 2100  C) showed that the smaller ZrB2 particles can be oxidized homogeneously and form also smaller ZrO2 particles resulting  in denser oxide scale  layer (Figs. 5 and 6).  It seems that in this case the residual porosity is not so dominant and finer microstructure with homogeneously distributed ZrB2 as well as SiC grains is more beneficial in oxidation protection. Results of weight gain measurements and surface analysis showed that the porosity higher than 7% led to the rapid oxidation of UHTCs. The weight gain of samples after oxidation at 1650  C shows different trends. While the weight gain of sample ZS2-1 (sintered at 2100  C) with  the highest density continuously  increases,  the weight gain of samples ZS2-2 and ZS2-3 remains nearly the same as at 1500  C. Sample ZS2-5 showed significant weight  loss during oxidation at 1650  C due to the evaporation of gaseous species, which was already observed in lower extent at 1500  C. The oxidation behaviour of ZrB2 -SiC composites beyond 1600  C was studied in-situ and also predicted by thermodynamic analysis [38,39]. Both studies showed that due to the CO evolution  from SiC oxidation, bubbles and voids are formed beneath the protecting oxide layer. This could explain massive weight  loss  in sample ZS2-5 with 10% of porosity. The temperature dependence of oxide scale thickness on two different sintering temperatures (1700  C and 1900  C) for samples ZrB2—20 wt.% SiC are shown in Fig. 4. The oxide scale thicknesses increased with the oxidation temperature as expected. The thickness of the oxide scale  in sample ZS2-2 with higher density was lower compared to sample ZS2-5 with 90.0% relative density. Based on these results the sintering temperature of 1900  C was used for the next experiments. The sintering  temperature affected  the density and  the grain size of ZrB2 and SiC in the final material, and thereby the oxidation resistance. It seems these two parameters play a crucial role in the improvement of oxidation resistance of studied ZrB2 -SiC UHTCs. During oxidation the smaller and homogeneously distributed ZrB2 particles allowed  the  formation of finer and denser ZrO2 layer at temperatures of 1100  C and 1300  C. On the other hand finer SiC particle provides more uniform distribution of protective glass phase at higher temperatures.  3.2.2.   Influence of SiC starting powder particle size  Oxidation resistance was evaluated according to mass changes and thickness of oxide scale. Mass changes of ZS1, ZS2-2 and ZS3 samples oxidized at different temperatures are shown in Fig. 7. Fig. 7 shows that weight gain of ZS3 sample  increased rapidly with the oxidation temperature and reached 34.1% at the highest temperature (1650  C). On the other hand samples ZS1 and ZS2-2 with nano-meter and sub-micrometer SiC grains showed smaller weight gain at all tested temperatures. As already stated in previous section, at lower temperatures the oxidation occurs preferably on the ZrB2 grains. The ZrB2 starting powder was the same in all samples. On the other hand the average particle size of three different starting SiC powders was different and played an important role in controlling the final density and microstructure of prepared UHTCs. This  led to the  inhomogeneous distribution of  large SiC grains  in ZrB2 matrix and formation of  local oxidation spots of ZrO2 (s) and  Fig. 7. Mass change vs. oxidation temperature of the tested samples, hot pressed at 1900   C and prepared from different SiC starting powders.  Fig. 8.  Light microscope image of silica glass formed by oxidation of big SiC grains in ZS3 sample after oxidation at 1500   C for 1 h.  In order to understand the effect of microstructural changes due to the sintering temperature on the oxidation resistance, samples ZS2-1, ZS2-2, ZS2-3 and ZS2-5 were subjected to oxidation up to 1650  C in static air atmosphere. Decreasing the sintering temperature reduced the grain size (Fig. 2) and refined the microstructure of  samples.  It  indicates  that finer microstructure may be more effective in protecting against oxidation of material. However the relative density of prepared UHTCs differs considerable and have to be also taken in account. Fig. 2 and Table 2 show that with the increasing sintering temperature the grain size of ZrB2 and SiC and the relative density of 1100  C UHTCs increase. It is known that at temperatures less than  the oxidation of Zr of Hf diborides is unaffected by the SiC additions and these diborides oxidize preferentially. This  leads to the oxide scale formation containing B2O3 , ZrO2 or HfO2 while SiC maintaines unoxidised. [37] Assuming the same chemical composition of ZS21, ZS2-2, ZS2-3 and ZS2-5 samples,  the weight of ZS2-5 sample increased rapidly after oxidation at 1100  C (Fig. 3). On the other hand  the smallest weight gain was measured  for sample ZS2-3 sintered at 1750  C followed by samples sintered at 1900  C (ZS22) and 2100  C  (ZS2-1). At  the first  look  the difference between samples ZS2-5 and ZS2-3 were inconspicuous, however mass gain for ZS2-5 sample was almost six times higher than for ZS2-3. The microstructure of both samples was  found to be highly homoge \\x0c', '3046   Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049  Table 3 Thickness of  the oxidation  layer of samples prepared  powders after oxidation at 1500   C.  from different SiC starting  Sample:   a) ZS1  b) ZS2-2  c) ZS3   Ox. layer thickness (\\u242em)  20 28 75  B2O3 (l) on the surface of oxidized ZS3 sample (Fig. 8). Therefore for all samples at 1100  C and 1300  C the changes  in the weight  were marginal, but became remarkable at higher temperatures. In the case of ZS1 and ZS2-2 samples with fine and homogeneously distributed SiC particles provided a steady supply of silica glass. In a case of inhomogeneous and random distribution of big SiC grains, as  in sample ZS3, the silica glass was formed only  locally and the oxidation could proceed quickly resulting in much lower oxidation resistance. The degradation of oxidation resistance was observed also by measuring the thickness of oxide scale after oxidation at 1500  C for 1 h (Fig. 9a-c, Table 3). The most excessive degradation of oxidation resistance was observed for sample ZS3 what is consistent with the results obtained by weight gain measurements. The emerging protecting glass layer is less compact and there is a deeper degradation of the material during oxidation. On the contrary, finer SiC particles in samples ZS1 and ZS2-2 caused the formation of thinner oxidation layer. Detailed view of the cross-sections of the samples exposed to oxidation is shown in Fig. 9. The layered structure with glass phase on the surface of oxidized samples was observed in all cases, which is consistent with available literature data [3,4,12-17]. As reported by Hwang et al. [22], uniform distribution of finer SiC grains  increases the ZrB2 /SiC  interface  length per unit area of exposed surface and thus decreases the spacing among SiC grains. This results in a faster formation of silica-rich layer on the surface of oxidized UHTCs. These considerations were  further proved by several studies [4,14,25] and revealed that the SiC grain-size reduction enhanced the oxidation resistance by the formation of thinner oxide scale.  3.2.3.   Influence of Y2O3 addition  The oxidation resistance of UHTCs depends on the formation of protective layer from the present compounds and oxidation products. The extent of protection  is determined by the physical and chemical processes as well as by the microstructure and composition of the oxidized material. Additives can also be used to improve the oxidation resistance of UHTCs. In the literature five main areas were  identified  for  the  improvement of  the oxidation resistance [3]:  (a)  Increasing the viscosity of borosilicate liquid, (b)  Inhibiting the polymorphic transformations of ZrO2, (c) Using alternatives to SiC to introduce silicon, (d) Formation of protective refractory phases at high temperature, (e) Modifying the microstructure of ZrO2 scale.  In this work the purpose of Y2O3 addition was to stabilize zirconia  formed during oxidation of ZrB2 and  to prove,  if Y2 Zr2O7 refractory phase will be  formed  at  the oxidation  temperature of 1650  C. RE2 Zr2O7 refractory zirconates have a high melting temperatures of above 2300  C and could provide  the oxidation protection at temperatures above 1600  C, when the borosilicate glass evaporates from the exposed surface [3]. The mass changes and thickness of oxidation  layers as a function of temperature for ZSY-2 and ZS2-2 samples (with and without addition of Y2O3 ) are  shown  in Figs. 10 and 11 Both  the mass  Fig. 9. SEM images of cross-sections of ZrB2 -SiC composites sintered at 1900   C and 30 MPa after oxidation at 1500   C, (a) ZS3 (44  \\u242em SiC), (b) ZS2-2 (0.9  \\u242em SiC), (c) ZS1 (55 nm SiC).  changes as well as thickness of layer formed after oxidation were higher in sample doped with Y2O3 . The most significant difference was observed after oxidation at 1650  C. The examination of the cross-sectioned surface after oxidation at 1500  C (Fig. 12) showed irregular oxidation layer with spallation and cracks. EDX analysis showed that on the top of the surface silica-rich glassy  layer was formed (Fig. 13a) and beneath this  layer aligned identified (Fig. 13b). At 1650  C the Y-doped ZrO2 particles were  glassy phase is invisible and only white oxide scale with thickness 630  of  \\u242em was observed (Fig. 14). XRD analysis showed that only monoclinic modification of zirconia was  found  in  the sample ZS2-2. Addition of Y2O3 caused  \\x0c', 'Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049   3047  Fig. 10. Plots depicting the mass changes depending on the temperature for ZS2-2 and ZSY-2.  Fig. 13. EDX analysis of ZSY-2 after oxidation at 1500   C; (a) Spot 1—silica-rich glassy layer on the top of surface (b) Spot 2—aligned Y-doped ZrO2 particles beneath silica glass layer.  Fig. 11. Plots of thickness layer formed during oxidation as a function of temperature for ZS2-2 and ZSY-2.  Fig. 12. SEM image of ZSY-2 after oxidation at 1500   C.  the formation of both, monoclinic and cubic phases of ZrO2 after oxidation at 1500  C as well as at 1650  C (Fig. 15). It was reported that the addition of Y2O3 improved both the densification and flexural strength of ZrB2 -SiC ceramics [31]. Guo et al. [33],  investigated  the addition of 3 vol.% RE2O3 rare-earth oxide  Fig. 14. Cross-section of sample ZSY-2 oxidized at 1650   C.  (RE = La, Nd, Y or Yb) to ZrB2 -20 vol.% SiC composites. They reported that Y2O3 and Yb2O3 addition enhances the densification of composite without ZrB2 grain growth and showed an  intergranular fracture mechanism resulting in a substantially increased hardness as well as fracture toughness. Moreover, RE2 Zr2O7 phase was found in Y2O3 and Yb2O3 doped composites located in direct contact with SiC grains. Recently, Zapata-Solvas et al. [35], showed that  in the ZrB2 -20 vol.% SiC doped with 2 wt.% of La2O3 oxide scale thickness is lower at 1400  C as for neat ZrB2 -20 vol.% SiC composite. On the other hand at 1500  C 1.4 times thicker and at 1600  C 2.5 times  \\x0c', '3048   Z. Ková ˇcová et al. / Journal of the European Ceramic Society 36 (2016) 3041-3049  Acknowledgements  This work was financially supported by the Austrian Ministry of Economic, Family and Youth within the frame of Research Studio Austria Programme  (RSA-RSP; 832018) of FFG. The authors’ acknowledge Dr. Christian Jogl, AAC GmbH, for SEM measurements.  References  Fig. 15. XRD patterns of ZSY-2 after oxidation at 1500   C and 1650   C.  indicating degradation of oxidation  thicker  layers were observed  resistance. Considering the massive weight gain and thickness of the oxide scale in ZSY-2 sample it seems to be clear that also Y2O3 , similarly as La2O3 causes degradation of oxidation  resistance at  temperthan 1500  C. The addition of Y2O3 atures higher  into ZrB2 -SiC composites proved the  formation of stabilized ZrO2 ; however no RE2 Zr2O7 phase was detected after oxidation at 1650  C. The oxidation resistance of Y2O3 doped ZrB2 -SiC sample (ZSY-2) was inferior in comparison to the basic yttria-free material (ZS2-2). Despite of  the  inferior  results  for  the oxidation  resistance of diboride-based UHTC with the addition of 8 wt.% Y2O3 , the formed crystalline oxide layer holds promise for improving their oxidation behaviour at higher temperatures (above 1700  C). However, additional research is needed to determine the mechanisms by which rare earth oxide additives affect  the structure of  the oxide scale along with the impact on transport of oxygen and/or gaseous oxidation products through the scale.  4. Conclusion  The effect of sintering temperature, SiC starting powder particle size and Y2O3 addition on the oxidation behaviour of ZrB2 -SiC ceramics were  investigated. The sintering temperature as well as average particle size of starting SiC powder played an  important role in controlling the final density and microstructure of prepared UHTCs, including controlling the ZrB2 grain size and thereby their temperatures of 1100  C and oxidation  resistance. At oxidation  1300  C, finer ZrB2 particles showed uniform oxidation resulting in homogeneous precipitation of very small ZrO2 particles  forming denser protecting  layer on  the surface of composite. On  the other hand, oxidation of finer SiC particles caused more uniform distribution of protective silica-based glass phase at higher temperatures. The most extensive oxidation of ZrB2 -SiC composite was observed in sample prepared from 44  \\u242em SiC powder. The obtained results give a hint that further improvement of oxidation resistance of ZrB2 -SiC UHTCs may be achieved by using of ZrB2 and SiC nanopowders. Addition of 8 wt.% Y2O3 to ZrB2 -20 wt.% SiC caused the formation of cubic ZrO2 , although no RE2 Zr2O7 phase was detected after oxidation at 1650  C. Nevertheless, the oxidation resistance of oxidation temperatures exceeding 1500  C. The weight gain as well Y2O3 doped samples was inferior in comparison to basic material at layer formed after oxidation at 1650  C was much as thickness of  higher for sample doped with Y2O3 as for the basic one. The thick630  \\u242em after oxidation at 1650  C ness of the oxide scale reached  for 1 h.  [2]   [5]   [15]   [12]   [11]   [1] E. 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},{
  "_id": 106,
  "PDF": "Influence of surface oxidation on the radiative properties of ZrB2-SiC composites.pdf",
  "Text": "['Applied Surface Science 409 (2017) 1-7  Contents lists available at ScienceDirect  Applied  Surface  Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / a p s u s c  Full Length Article  Influence of surface oxidation on the radiative properties of ZrB2-SiC composites  Ning Li a,∗ , Pifeng Xing a , Cui Li a , Peng Wang b , Xinxin Jin c , Xinghong Zhang d  a Research Center of Laser Fusion, China Academy of Engineering Physics, Mianyang, 621900, PR China b School of Material Science and Engineering, Shandong University of Technology, Zibo 255049, PR China c College of Materials Science and Engineering, Harbin University of Science and Technology, Harbin 150040, PR China d Science and Technology on Advanced Composites in Special Environments Laboratory, Harbin Institute of Technology, Harbin 150001, PR China  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 22 January 2017 Received in revised form 17 February 2017 Accepted 28 February 2017 Available online 3 March 2017  Keywords:  ZrB2 -SiC Emissivity Radiative properties Surface oxidation  1.   Introduction  The spectral emissivities of ZrB2 -20 vol.% SiC composites with various surface components of ZrB2 /SiC (ZS1), silica-rich glass (ZS2) and porous zirconia (ZS3) were measured using  infrared spectrometer  in the wavelength range from 2.5 to 25.0  \\u242em. The relationship between surface oxidation (associated with surface component, thickness of oxide scale, testing temperature as well as roughness) and the radiative properties of ZrB2 -SiC composites were  investigated systematically. Surface component affected the radiative properties of composites significantly. The total emissivity of ZS1 varied from 0.22 to 0.81 accompanied with surface oxidation  in the temperature range 300-900   C. The emissivity of ZS2 was about 1.5 times as that of ZS3 under the same testing conditions. The oxide scale on specimen surface enhanced the radiative properties especially in terms of short-wave range, and the emissivity in the longwave range gradually increased with the thickness of oxide scale within a certain range. The influence of testing temperature and surface roughness was also investigated. The testing temperature had a little effect on radiative properties, whereas effect of surface roughness could be negligible. © 2017 Elsevier B.V. All rights reserved.  Aerospace materials research in the last decade has been focused on ultra-high temperature ceramics (UHTCs)  for a variety of elevated temperature structural applications, such as nose tips and sharp wing  leading edges on hypersonic vehicle  [1-3]. As one important family of UHTCs, zirconium diboride (ZrB2 ) based composites have some advantages over other materials, including high retained  strength and good  resistance  to oxidation along with relatively low density, and are being considered to operate at temperature above 1600  C  in severe environments  for a  long  time [4-7]. However, high surface temperature and aerodynamic heating will be detrimental  to  the high-temperature performance of composites, which can result  in unexpected catastrophic  failure of thermal protection materials during hypersonic re-entry flight. Generally, there are two ways available to reduce the materials’ surface temperatures under re-entry condition. One method involving surface catalytic effect  is to  inhibit the transfer of recombination energy  from dissociated atoms to material surface  [8,9]. For this  ∗ Corresponding author. E-mail address: lncaep@163.com (N. Li).  http://dx.doi.org/10.1016/j.apsusc.2017.02.266 0169-4332/© 2017 Elsevier B.V. All rights reserved.  aim, SiC and silica glass with  low catalytic coefficients have been used as coatings [10,11]. The other approach, which  is related to this study, is to increase the emissivity of thermal protection material. Because the heat flux of wall material at low pressure is mainly dissipated by radiation. Radiative properties are pivotal to thermal protection materials as candidates  for aerospace applications, and emissivity  is also a key parameter for computational fluid dynamics (CFD) simulations. Numerous experiments have been  reported on  the emissivities of materials under various testing conditions. The effects of temperature on  radiative properties of different  thermal protection materials  for reusable space vehicles were reported  [12,13]. The emissivity of Y2O3 rose  in  the  temperature  range 750-1800 K, whereas  the emissivity of HfO2 stayed nearly constant and  the emissivity of Al2O3 decreased with increasing temperatures. Zhou et al. [14] demonstrated the relationship between emissivity and thickness of Al2O3 coating, indicating that the emissivity of coating 0.9 above thickness of increased with thickness and stabilized at  40  \\u242em. Additionally, high sensitivity of radiative behaviors to surface roughness was also confirmed by Auweter-Kurtz et al. [12] for SSiC, Rousseau et al. for Pr2NiO4 coating [15] and Wen et al. for aluminum alloys  [16]. However, the emissivity  for some nonmetals particularly white ceramic materials appeared to be  independent                                        \\x0c', '2   N. Li et al. / Applied Surface Science 409 (2017) 1-7  of surface  roughness, at  least  for wavelengths below 7 or 8  \\u242em [17]. For ZrB2 -based composites, the specimens prepared by electrical discharge machining showed higher emissivity than the ones machined by diamond-loaded tools at temperatures below 1500 K, but  the effect of surface finish became negligible at higher  temperature [18]. Tan et al. [19] reported that incorporating rare-earth oxide Sm2O3 into ZrB2 -SiC coating enhanced the radiative properties, and  the emissivity declined first and  then gradually  rose as  the  temperature  increased. Recently,  influence of SiC  shape on the  infrared emissivity properties has been  investigated  [20]. Compared with SiC particles, the addition of SiC whiskers as the second phase  into ZrB2 was beneficial  to  radiative properties. Radiative behaviors of ZrB2 -based composites are quite complicated. Meanwhile, ZrB2 -based composites are usually exposed to high-temperature environment and accompany surface oxidation during  service. The  relationship between  surface oxidation and emissivity has not been studied sufficiently, thus a comprehensive consideration on radiative behaviors  is necessary  for ZrB2 -based composites, especially in oxidizing environment. In  the present work, ZrB2 -20 vol.% SiC composites were preoxidized at 1500  C for 30 min  in static air and at 1000 Pa  in pure oxygen, respectively. The spectral emissivities of composites were \\u242em. The measured using  infrared spectrometer  from 2.5  to 25.0  purpose of  this paper  is  to  investigate  the  radiative properties of original ZrB2 -SiC specimen and pre-oxidized specimens  in the temperature range 300-900  C. The influence of surface oxidation, associated with surface component, thickness of oxide scale, testing temperature and roughness, on the emissivity was discussed in detail.  It  is believed that this work could  form a  further comprehension to the evolution mechanism of radiative behavior for ZrB2 -based composites, and provide data  for numerical simulations.  Fig 1. Schematic presentation of the emissivity test apparatus.  spectral radiation intensity between blackbody and specimen at a given temperature, which can be represented as  ε((cid:3), T )   =  E(cid:3) ((cid:3), T ) E(cid:3),b ((cid:3), T )  (1)  where ε((cid:3), T)  is spectral emissivity at the temperature T; E(cid:3) ((cid:3), T) and E(cid:3),b (\\u242d, T) are spectral radiation of specimen and blackbody, respectively.   denotes the wavelength. In addition, the total emissivity ε(T) of specimen from 2.5 to 25.0  \\u242em can be calculated by the following equation:  \\u242d  (cid:2) 25  2.5 ε((cid:3), T )E(cid:3),b ((cid:3), T )d(cid:3)  (cid:2) 25  2.5 E(cid:3),b ((cid:3), T )d(cid:3)  2. Experimental procedure  =  ε(T )   (2)  ×  Commercially  available  ZrB2 powder  (d50 = 2.1  \\u242em,  purity >99.5%, Northwest Institute for non-ferrous metal research, China), and  SiC powder  (d50 = 1.4  \\u242em, purity  >99.5%, Weifang Kaihua Micro-powder  Co.  Ltd.,  China) were  used  as  starting materials. We have described  the preparation procedure of ZrB2 -20 vol.% SiC composites  in detail previously  [21]. Bulk density and theoretical density were measured by  the Archimedes method following ASTM C373-88 and  the  rule-of-mixture,  respectively. Round-shaped specimens with Ø30 mm   3 mm were cut from the billet, and polished with diamond slurries. The ZrB2 -SiC original specimen is designated as ZS1 for simplicity. The pre-oxidation was carried out in static air using a furnace and at 1000 Pa in pure oxygen using a side-arm oxidation facility [22] at a constant temperature of 1500  C for 30 min, the corresponding as-treated specimens are named as ZS2 and ZS3, respectively. Infrared emissivity test apparatus, as  illustrated  in Fig. 1, was used to measure the normal spectral emissivity of ZrB2 -SiC composites. Specimens were heated in static air up to 300, 500, 700 and 900  C, respectively. The surface temperature was monitored by a K-type thermocouple, and the blackbody furnace (P1200B, Landcal, UK) as the near-blackbody source was held at the same temperature. Temperatures of specimen and blackbody were managed by a temperature controller and a compensation system during testing. Fourier transform infrared spectrometer (FTIR-6100, Jasco, Japan) was applied to detect  infrared radiance with a range  from 2.5 to 25.0  \\u242em. Background noise was removed based on  the environmental radiation compensation algorithm to eliminate disturbance. The spectral emissivity of specimen is obtained by comparing the  The  surface microstructural  features  of  specimens were observed by scanning electron microscopy (SEM, FEI Sirion, Holland) equipped with energy dispersive spectroscopy (EDS, EDAX Inc., USA) for chemical analysis. Crystalline phases of ZS1 after testing were analyzed by X-ray diffraction (XRD, Rigaku,  Japan) that was operated at an incidence angel of 2  with a 2\\u242a  range from 10  to 80  . The surface chemical composition was identified using X-ray photoelectron spectroscopy  (XPS, Thermo Fisher Scientific, USA) with monochromatic Al K␣  radiation. Optical 3D profiler (Sensofar, Spain) was used to measure the surface roughness.  3. Results and discussion  Photographs of ZrB2 -20 vol.% SiC specimens are shown in Fig. 2. Surface of original ZrB2 -SiC specimen  (ZS1) after polishing was quite flat and bright  (Fig. 2a). The  surface colors of  specimens with different pre-oxidation conditions, as shown  in Fig. 2b and c, changed into dark and white for ZS2 and ZS3, respectively. This difference of surface morphology can be mainly attributed to the variation of specimen surface associated with chemical composition as well as microstructure. The in situ measurement of emissivities for ZrB2 -SiC specimens was performed using  infrared emissivity test apparatus.  In Fig. 3 spectral emissivities are plotted versus temperature  for different batches of ZrB2 -SiC specimens.  It was evident  that  the variation of the spectral emissivity of ZS1 as a  function of testing temperature was significantly different from those of ZS2 and ZS3. After  \\x0c', 'N. Li et al. / Applied Surface Science 409 (2017) 1-7   3  Fig. 2. Photographs of ZrB2 -20 vol.% SiC specimens: (a) ZS1 (original specimen), (b) ZS2 (specimen pre-oxidized in air) and (c) ZS3 (specimen pre-oxidized at 1000 Pa).  Fig. 3. Spectral emissivity of ZrB2 -20 vol.% SiC specimens as a function of wavelength: (a) ZS1, (b) ZS2 and (c) ZS3.  Table 1 Summary of measured emissivity, roughness and thickness of oxide scale for ZrB2 -20 vol.% SiC specimens.  Specimens   ZS1  ZS2  ZS3  Testing temperature (  C)   Roughness (nm)   Thickness of oxide scale (\\u242em)   Emissivity  300  500  700  900   300  500  700  900   300  500  700  900   41.7  59.2  148.6  73.5   430.1  509.6  385.8  416.3   3956.6  2423.8  2629.5  1905.7   0  0  0.8  4.1   27.4  25.7  27.2  29.2   58.3  60.7  55.8  57.4   0.22 0.26 0.56 0.81  0.74 0.76 0.83 0.85  0.44 0.49 0.52 0.54  testing at 300 and 500  C, the spectral emissivities of ZS1 were relatively low, varied between 0.61 and 0.13 in the wavelength range from 2.5 to 25.0  \\u242em. As the testing temperature rose, a substantial  increase  in spectral emissivity was observed after  testing at 700  C, and the measured value at 900  C  increased further espe\\u242em. From Fig. 3b, there was a cially in the long wave range above 10  slight difference in spectral emissivity over the entire testing temperature range as presented, and emissivity curves of ZS2 tested  at different  temperatures were similar. Compared with ZS2, ZS3 spectral emissivities were quite low, grow from 0.01 to 0.63 with the wavelength (Fig. 3c). All the ZrB2 -SiC specimens’ total emissivities, as summarized  in Table 1,  increased as testing temperature rose. Nevertheless, ZS2 and ZS3 had low sensitivity of emissivity to temperature which varied from 300 to 900  C. ZS2 exhibited a minimum value of emissivity of 0.74 at 300  C and a maximum value of 0.85 at 900  C, meanwhile the measured values of ZS3 ranged from  \\x0c', '4   N. Li et al. / Applied Surface Science 409 (2017) 1-7  0.44 to 0.54. On the contrary, ZS1 had an emissivity of 0.22 at low temperature, rising to 0.81 with the testing temperature of 900  C. It is believed that the evolution of emissivity is strongly dependent on surface composition and morphology of ZrB2 -SiC specimens. The surface microstructures of three batches of ZrB2 -SiC specimens are shown in Fig. 4. For the original ZrB2 -SiC specimen (ZS1), 99% using the rule mixtures based the calculated bulk density was  on the densities of 6.09 and 3.21 g/cm3 for ZrB2 and SiC, respectively [23]. Except for small dark SiC dispersed uniformly in the gray ZrB2 matrix, no pits, pores or flaws can be identified from Fig. 4a. Fig. 4b and c display as-treated specimens’ surface morphologies, which were essentially different from that of ZS1. The surface of ZS2 was dark and completely covered with a smooth and continuous glass, resulting from the pre-oxidation at 1500  C in static air by Eqs. (3)(5). Combined with EDS analysis, as shown  in Fig. 4d, the porous surface of ZS3 was mainly composed of zirconia. Comparison of the surface morphologies for specimens before and after pre-oxidation at 1000 Pa  in pure oxygen (Fig. 4a and c), distribution of ZrO2 in ZS3 surface was consistent with that of ZrB2 in matrix. The pores in  the surface can be ascribed  to  the depletion of SiC according to Eq. (6), while a ZrB2 grain was oxidized  in situ to several ZrO2 grains (Eq. (7)). It can be expected that transportation, agglomeration and growth of ZrO2 with liquid products could be neglected owing to the active oxidation of SiC associated with  low oxygen partial pressure [24].  →  →  +   B2O3 (l)   (3)  (4)  (5)  (6)  (7)  +  +  +  +  ZrB2 (s)    5/2O2 (g)    ZrO2 (s)   SiC(s)    3/2O2 (g)   xB2O3 (l)    ySiO2 (l)   +  +   SiO2 (l)   CO(g)   xB2O3 ·  ySiO2 (l)   →  SiC(s)    O2 (g)    SiO(g)    CO(g)   →  +  ZrB2 (s)    5/2O2 (g)   →   ZrO2 (s)   +   B2O3 (g )   The surface SEM images and XRD patterns of ZS1 after testing, as shown in Fig. 5, provide insight into the evolution of microstructure and condensed phase of specimens oxidized  in  the  temperature range 300-900  C. It should be noted that the difference in macroscopic appearance as well as microstructure could be neglected the emissivity of ZS1 at 300 and 500  C. From after measuring  Fig. 5a, the surface of ZS1 tested at 500  C was mainly composed of ZrB2 and SiC, and a few small pores on surface could be attributed to the grains being pulled out during polishing and/or being oxidized during  testing. However, TGA analysis  reported by other investigators [25-27] has demonstrated that the  initial oxidation temperatures were 700  C  for ZrB2 particle, 900  C  for SiC particle, respectively. Fig. 5b is micrograph of ZS1 surface after testing at 700  C. The surface was covered with thin B2O3 glass, but ZrB2 and SiC grains under outer glass  layer could be  identified, which the specimen was oxidized slightly at 700  C. As indicated  that  the testing temperature further increased to 900  C, no other condensed phase except for B2O3 glass was observed (Fig. 5c). It was apparent that the thickness of outer glass increased with the testing temperature. Fig. 5d displays the XRD patterns of ZS1 tested at different temperatures. The surfaces of specimens tested at 300 and 500  C were composed of ZrB2 as well as a small quantity SiC, no peaks corresponding to other phases could be detected. With rising testing temperature, m-ZrO2 appeared  at 700  C, and then became the predominant component accomin the surface scale panied with disappearance of peaks for ZrB2 and SiC at 900  C. For ZS2 and ZS3, the amorphous borosilicate glass and zirconia skeleton on surfaces after pre-oxidation can act as protective barriers to inhibit the diffusion of oxygen into the inner part of the matrix. The oxidation of pre-oxidized specimens during the testing of radiative properties could be negligible, and the surface components as well as morphologies of ZS2 and ZS3 are not  fundamentally different from those after testing in the temperature range 300-900  C.  XPS was also performed  to  identify  the evolution of surface chemical composition  for ZS1 after  testing at various  temperatures. The specimen surfaces were cleaned by Ar+  ion sputtering for 1 min, and C 1s signal detected at 285.0 eV  from adventitious carbon was utilized to calibrate the spectra. As shown  in Fig. 6a, Zr  B bond as the main component could be observed at 178.7 eV (Zr 3d5/2 ) and 181.1 eV (Zr 3d3/2 ) from Zr 3d spectra after testing at 300  C, while two weak peaks located at 182.8 eV (Zr 3d5/2 ) and 185.1 eV (Zr 3d3/2 ) corresponded to Zr  O bond.  It was suggested that  there was some slight oxidation  for ZrB2 at 300  C. As  the testing temperature grew, the characteristic peaks of Zr  B bond in spectra became weak markedly; meanwhile, the relative intenO bond increased first at 500  C, then gradually declined sity of Zr  and disappeared from Zr 3d spectra at 900  C. Fig. 6b presents the development of curve fitting Si 2p spectra for ZS1 in the temperature range 300-900  C. After testing at 300  C, the main component located at 100.6 eV was confirmed as SiC. Aside  from Si  C bond, Si O bond can be detected at 103.2 eV when the testing temperature increased. It was believed that SiC has been oxidized mildly at 500  C. Similarly,  the  intensity of Si  O bond  increased as  the temperature further went up to 900  C. From Fig. 6c, the chemical shift between O  Zr bond (530.8 eV) and O-B bond (533.2 eV) could be identified in O 1 s spectra associated with the specimen surfaces after testing. It was evident that ZrO2 was the main oxidation product on the surface at 300 and 500  C. When the temperature rose to 700  C, O  B bond became the main component, which was  in relation to B2O3 glass covering the surface (as indicated in Fig. 2c). Compared with ZS1 tested at 700  C, no detectable change in specZr bond was detected at 900  C. The disappearance tra except for O  of O  Zr signal indicated that B2O3 glass layer became much thicker, and resulting ZrO2 was completely covered after testing. Comparison among emissivities  for different batches of ZrB2 SiC specimens suggested that the surface component significantly affected, as expected, the radiative properties of composites. The emissivity of ZS2 was about 1.5 times as that of ZS3 under the same test condition, because resulting borosilicate glass on the surface of ZS2 showed high  radiative properties compared with porous ZrO2 on ZS3 surface. Similarly, change  in emissivity of ZS1 could be explained by  the evolution of surface component at various testing temperatures. After testing at 300 and 500  C, XPS analysis  indicated that trace oxygen existed on ZS1 surface, but  it was reasonable to consider that the emissivity of ZrB2 -SiC composites without oxidation was quite  low, even below that of ZrO2 at  low to 700  C,  temperature. As  the  temperature grew  from 500  the emissivity rose by 115.4% that was attributed to the variation of surface component caused by slight oxidation. Furthermore, when the temperature rose further to 900  C, the emissivity of ZS1 was almost same with that of ZS2. Based on  the  above-mentioned discussion,  the  variation of surface component associated with oxidation during testing was responsible for the remarkable increase in emissivity of ZS1 tested above 700  C. However, an apparent difference in emissivity curves tested at 700 and 900  C. As  was observed  from ZS1  indicated in Fig. 3 and Table 1,  the emissivity of ZS1  increased mainly  in short-wave range at 700  C and long-wave range at 900  C, which could be related to the thickness of the oxide scale. After testing at 700  C, the thickness of oxide scale on surface was relatively thin as only 0.8  \\u242em. It was known that long-wave of matrix could penetrate oxide scale compared with short-wave. Therefore, existence of oxide scale on the surface had great  influence on the radiative properties in terms of short-wave range. As the testing temperature rose to 900  C, almost of radiance detected by  infrared spectrometer came from oxide scale, and the effect of surface oxidation on emissivity in long-wave range gradually increased. That  the radiative properties of  thermal protection materials are associated with surface  temperature has been confirmed by  \\x0c', 'N. Li et al. / Applied Surface Science 409 (2017) 1-7   5  Fig. 4. Micrographs and EDS analysis of the surface for ZrB2 -20 vol.% SiC specimens: (a) ZS1, (b) ZS2 and (c) ZS3 and (d) EDS spectra of (c).  Fig. 5. Micrographs and XRD patterns of ZS1 after testing: (a) 500   C, (b) 700   C, (c) 900   C and (d) XRD patterns.  some  literatures  [12,13,28];  indeed,  the emissivities of ZS2 and ZS3, as  indicated  in Fig. 3 and Table 1,  rose  in  the  temperature to 900  C. For ZS1,  range  from 300  it was expected  that  testing temperature would contribute  to  the change  in emissivity simi larly, however, the impact of temperature on radiative properties of ZrB2 -SiC specimens was limited compared with the effect of the surface component.  \\x0c', '6   N. Li et al. / Applied Surface Science 409 (2017) 1-7  Fig. 6. XPS spectra of ZS1 after testing at different temperatures: (a) Zr 3d, (b) Si 2p and (c) O 1s.  Early experiments by Rousseau et al. and Wen et al. showed that a rough surface enhanced the emissivity [15,16]. The roughness of mirror-like surface of ZS1 tested at 500  C was rather small, which was beneficial to  low radiative properties. However, when the surface roughness decreased from 148.6 to 73.5 nm, there was no detectable effect of  roughness on emissivity with  increasing to 900  C. Similarly,  temperature  from 700  in spite of  the  large roughness for ZS3, the emissivity of ZS3 was much lower than that of ZS2. Therefore, compared with other factors as mentioned above, surface roughness had  little effect on  the radiative properties of ZrB2 -SiC composites.  4. Conclusion  ZrB2 -20 vol.% SiC composites pre-oxidized under different conditions  (static air and 1000 Pa pure oxygen) were covered with silica-rich glass and zirconia, respectively. The spectral emissivities of ZrB2 -SiC specimens with various surface morphologies were measured using  infrared spectrometer  in  the  temperature range 300-900  C, in order to investigate the influence of surface oxidation on the radiative properties.  ZrB2 -SiC composites’ emissivities were significantly affected by surface component. The emissivity of ZS1 rose by 115.4% accompanied with slight oxidation as the temperature grew from 500 to 700  C. Similarly, the emissivity of ZS2 related to resulting borosilicate glass on surface was about 1.5 times as that of ZrO2 on ZS3 surface under the same testing condition. The variation of radiative properties for ZS1 tested at 700 and 900  C was mainly attributed to the different thicknesses of oxide scales. The effects of surface oxidation on emissivity in long-wave range gradually increased with the thickness of oxide scale within a certain range. The emissivities of composites increased as the testing temperature rose from 300 to 900  C, indicating that testing temperature was also responsible for the change in radiative properties, but the impact was limited as compared to surface component. The radiative properties of ZrB2 SiC composites were not sensitive to surface roughness compared with other factors as mentioned above. High-temperature performance of thermal protection materials is closely related to surface temperature, which is strongly dependent on  the emissivity. For  the ZrB2 -based composites, covering the surface with borosilicate glass by pre-oxidation in static air is a promising method for improving radiative properties of composites.  \\x0c', 'N. Li et al. / Applied Surface Science 409 (2017) 1-7   7  Acknowledgments  This work was financially supported by  the National Science Foundation of China (Project Nos. 51272056, 11121061, 51372047 and 91216301), the National Fund for Distinguished Young Scholars (No. 51525201) and the Fundamental Research Funds  for the Central Universities (Grant No. HIT.BRET III. 201506).  References  [5]   [1] D.M. Van Wie, D.G. Drewry, D.E. King, C.M. Hudson, The hypersonic environment: required operating conditions and design challenges, J. Mater. Sci. 39 (2004) 5915-5924. [2] E. Wuchina, E. Opila, M. Opeka, W. Fahrenholtz, I. Talmy, UHTCs Ultra-high temperature ceramic materials for extreme environment applications, Electrochem. Soc. Interface 16 (2007) 30-36. [3] R. Savino, M.D.S. 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Meng, Oxidation behavior of zirconium diboride-silicon carbide at 1800   C, Scripta Mater. 57 (2007) 825-828. [25] W.M. Guo, G.J. Zhang, Y.M. Kan, P.L. Wang, Oxidation of ZrB2 powder in the temperature range of 650-800   C, J. Alloys Comp. 471 (2009) 502-506. [26] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [27] Z.J. Wu, Z. Wang, G.D. Shi, J. Sheng, Effect of surface oxidation on thermal shock resistance of the ZrB2 -SiC-ZrC ceramic, Compos. Sci. Technol. 71 (2011) 1501-1506. [28] H.F. Clark, D.G. Moore, A rotating cylinder method for measuring normal spectral emittance of ceramic oxide specimensfrom 1200 to 1600 K, J. Res. NBS 70A (1966) 393-415.  [17]   [24]   \\x0c']"
},{
  "_id": 107,
  "PDF": "Influence-of-YO-addition-on-the-mechanical-and-oxidation-behaviour-of-carbon-fibre-reinforced-ZrBSiC-composites2020Journal-of-the-European-Ceramic-Society.pdf",
  "Text": "['Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e lsev ie r .com / loca te / jeu rce ramso c  Original Article  Influence of Y2O3 addition on the mechanical and oxidation behaviour of carbon fibre reinforced ZrB2/SiC composites  Antonio Vinci*, Luca Zoli, Pietro Galizia, Diletta Sciti  CNR-ISTEC,  Institute of Science and Technology for Ceramics, Via Granarolo 64,  I-48018 Faenza,  Italy  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ceramic-Matrix Composites (CMCs) Ultra-High-Temperature-Ceramics (UHTCs) Fibre-matrix interface Rare earths Oxidation resistance  The influence of Y2O3 addition on the microstructure, thermo-mechanical properties and oxidation resistance of carbon fibre reinforced ZrB2/SiC composites was investigated. Y2O3 reacted with oxide impurities present on the surface of ZrB2 and SiC grains and formed a liquid phase, effectively lowering the sintering temperature and allowing to reach full density at 1900 °C. The presence of a carbon source (fibres) led to additional reactions which resulted in the formation of new secondary phases such as yttrium boro-carbides. Mechanical properties were significantly enhanced compared to the un-doped composite. Further tests at high temperatures resulted in strength increase up to 700 MPa at 1500 °C which was attributed to stress relaxation. Oxidation tests carried out at 1500 °C and 1650 °C in air showed that the presence of the Y-based secondary phases enhanced the growth of ZrO2 grains, but offered limited protection to oxygen due to the lower availability of surficial SiO2 formed from SiC.  1.  Introduction  The demand for materials able to withstand more challenging and harsh conditions than those encountered during re-entry in earth atmosphere or hypersonic flight, and surpass the current limits of C/SiC based CMCs, has driven researchers towards the study of a novel class of refractory ceramics called ultra-high temperature ceramics (UHTCs). These comprise the carbides and borides of early transition metals which are characterized by melting points above 3000 °C and are being considered as candidates for the application in extreme environments. Amongst UHTCs, ZrB2 is the most commonly studied due to its relatively low density, high thermal conductivity, and lower price compared to HfB2 [1,2]. The main drawback is the low fracture toughness of these materials that limits their application where thermal shocks and vibrations are present [3]. Moreover the oxidation resistance of pure ZrB2 is poor because of the formation of a porous and non-protective ZrO2 scale and the evaporation of B2O3 already at 1000 °C. In this regard, many efforts were made to overcome the low oxidation resistance of ZrB2 by introducing additives, such as SiC and other silicides, which promote the formation of a surficial glassy silica phase which protects the material from further oxidation up to 1600 °C. In order to improve the damage tolerance of these materials, short and long carbon fibres have been extensively investigated as reinforcement [4-14]. Results show how the introduction of the fibre reinforcement  lowers the overall strength of the composite but greatly improves the damage tolerance and thermal shock resistance. Recently these materials were validated in an arc-jet wind tunnel facility and are currently being scaled up [15,16]. The main issues encountered during the fabrication of fibre reinforced UHTCs are typically related to the problematic infiltration of the fibre preforms and the consolidation of the green composite. Studies on the sintering behaviour of these composites showed how temperatures above 1900 °C allowed to reach higher densities but severely damaged fibres in the process, therefore the addition of sintering aids was further investigated. As far as the oxidation resistance is concerned, an oxidation kinetics study carried out on CfZrB2/SiC composite showed how below 1100 °C only the fibres and ZrB2 are oxidized, leading to the formation of B2O3 which is non-protective and tends to evaporate [17]. At higher temperatures, the rapid formation of SiO2 is triggered, leading to the passivation of the material up to 1550 °C. Since these materials are designed to operate at very high temperatures that often involve high heat fluxes, short term oxidation tests often provide a good starting point to evaluate the potential of these materials to withstand extreme environments. A study on the short term-oxidation resistance of these materials showed how exposing these materials directly to temperatures above 1500 °C allows to bypass the critical stage mentioned before while providing a tool to compare the oxidation behaviour [14]. Rare earth oxides such as Y2O3 have been shown to improve the  ⁎ Corresponding author. E-mail addresses: antonio.vinci@istec.cnr.it (A. Vinci),  luca.zoli@istec.cnr.it (L. Zoli), pietro.galizia@istec.cnr.it (P. Galizia), diletta.sciti@istec.cnr.it (D. Sciti).  https://doi.org/10.1016/j.jeurceramsoc.2020.06.043 Received 27 February 2020; Received in revised form 8 June 2020; Accepted 13 June 2020 0955-2219/ © 2020 Elsevier Ltd. All rights reserved.  Please cite this article as: Antonio Vinci, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2020.06.043  \\x0c', 'A. Vinci, et al.  densification of ZrB2/SiC bulk composites and allow to reach near full density ceramics due to the formation of a liquid phase with the oxide impurities present on the surface of ZrB2 and SiC grains [18,19]. Moreover the presence of Y2O3 is thought to stabilize the ZrO2 formed during oxidation in the tetragonal structure, leading to the formation of a more compact scale [20]. However, the effect of Rare Earth oxides on complex systems such as fibre reinforced UHTCs, where carbon is also present, has been never investigated. In this work, 5 vol% Y2O3 was added to Cf reinforced ZrB2-SiC composites in order to improve densification, mechanical behaviour and oxidation resistance. Short term oxidation tests were carried out to explore the potential of Y2O3 for the oxidation resistance. The temperature ranges selected of 1500 and 1650 °C are based on previous works carried out on fibre reinforced ZrB2/SiC composites. The effect of Y2O3 on the high temperature mechanical behaviour up to 1500 °C was investigated.  2. Experimental  2.1. Raw materials  For the preparation of the ceramic suspensions, raw powders available commercially were used. The raw materials used were: ZrB2 (H.C. Starck, grade B, Germany, specific surface area 1.0 m2/g, particle size range 0.5−6 μm, impurities (wt.%): 0.25 C, 2 O, 0.25 N, 0.1 Fe, 0.2 Hf), α-SiC (H.C. Starck, Grade UF-25, specific surface area 23-26 m2/g, D50 0.45 μm, Italian retailer: Metalchimica), Y2O3 (H.C. Starck, 99.5 %, Grade C, specific surface area 10−16 m2/g, D50 0.90 μm, impurities (wt.%): 0.005 Al, 0.003 Ca, 0.005 Fe). Unidirectional high modulus carbon fibres (Granoch, Yarn XN80-6 K fibres; E =780 GPa, σtensile =3.4 GPa, Ø =10 μm) were used as fibre reinforcement.  2.2. Processing  Two powder mixtures containing ZrB2 + 5 vol% SiC (designated ZS) and ZrB2 + 5 vol% SiC + 5 vol% Y2O3 (designated ZSY) were prepared by wet ball milling of the commercial powders for 24 h and then dried with a rotary evaporator. The composites were fabricated via slurry infiltration of the fibre bundles; the fibre layers were piled up in a 0/0° configuration and then cut into a 30 × 30 mm pellet [21]. The sample was hot pressed at 1900 °C for both composites, using a pressure of 40 MPa and a holding time of 15 min, in accordance to previous studies [21].  2.3. Microstructure analysis  The microstructures of the sintered materials were analysed on polished and fractured surfaces with a field emission scanning electron microscope (FE-SEM, Carl Zeiss Sigma NTS Gmbh Öberkochen, Germany) and energy dispersive X-ray spectroscopy (EDX, INCA Energy 300, Oxford instruments, UK). X-ray diffraction analysis (Bruker D8 Advance apparatus, Karlsruhe, Germany) was carried out on the materials before and after oxidation tests. The fibre, matrix and phases volumetric amounts were evaluated by image analysis with software Image-Pro Analyser 7.0 and the theoretical densities of the composites were calculated using the rule of mixtures. Grain size was determined  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  through image analysis using standard ISO/DIS 13383-1, method A2.  2.4. Mechanical  testing  Flexural strength was measured by four-point bending on specimens with size 25 × 2.5 × 2 mm3 (Length × Width × Height) using a fullyarticulated steel fixture and a screw-driven testing machine (Zwick/ Roell, model Z050). The lower and upper span were 20 mm and 10 mm respectively, while the cross-head rate was 1 mm/min. The tests were carried out following the guidelines of standard EN 843-1 (2006). For the tests at 1200 °C and 1500 °C, a screw-driven testing machine (1195, INSTRON) was used; the specimens were placed on a semi-articulated alumina 4-point fixture and heated up to 1500 °C with a rate of 10 °C/ min under argon flow (3.5 L/min) in a high temperature furnace (HTTF model 924, Severn Furnaces Limited). Bars were held at 1500 °C for 15 min before testing. Due to inaccuracy of the apparatus in measuring the true strain of the load-displacement, the true strength of the material was considered reliable up to the proportional limit in the curve. This was calculated using the maximum load in the elastic region which was determined from the best fit of the linear portion of the load/displacement curves which had R2 values of 0.998. The fracture toughness (KIc) was evaluated by 4-point bending on chevron notched beams (CNB), following the guidelines of EN 14425-3 (2010). The equation of Munz et al. was used to calculate KIc [22]. The test bars, 25 × 2 × 2.5 mm3 (Length × Width × Height), were notched with a 0.1 mm-thick diamond saw; the chevron-notch tip depth and average side length were about 0.12 and 0.80 of the bar thickness, respectively. The testing apparatus is the same used for RT flexural strength. A crosshead speed of 0.05 mm/min was used.  2.5. Short-term oxidation tests  Bars with dimensions 2.5 × 2 × 12 mm were machined from the sintered specimens. The samples were cleaned with ethanol and dried under IR light. A preliminary thermo-gravimetric analysis was carried out on the doped and undoped specimens from 25 °C to 1500 °C with a heating rate of The oxidation tests were carried out in a bottom-up loading furnace (FC18-0311281, Nannetti Antonio Sauro S.R.L., Italy) at 1500 and 1650 °C in air for 1 min following the same procedure reported in [14]. The furnace was heated to the desired temperature with a rate of 5 °C/min. Then the specimens were placed in the furnace using a porous zirconia sample holder. After reaching the target temperature, the samples were held in the furnace for 1 min. At the end of the oxidation test, the specimens were quickly taken out and cooled down naturally in air.  3. Results and discussion  3.1. Microstructure of  the sintered material  The physical properties of samples ZS and ZSY are reported in Table 1. The final densities ranged from 3.7 to 4 g/cm3, depending on the amount of residual porosity and fibre volumetric content. The composites contained comparable amounts of fibres; slight deviations are due to the intrinsic variability and scatter of the manual process of infiltration. Since some chemical reaction took place during sintering,  Table 1 Volumetric amounts of the phases measured by image analysis and detected by XRD after sintering, density values (theoretical, experimental and relative), porosity and ZrB2 grain size for samples ZS and ZSY.  Sample  ZS ZSY  Composition (vol%)  ZrB2  49.41 45.76  SiC  2.64 1.10  ZrC  0.05 0.25  YB2C2  - 2.49  Cf  47.9 50.4  ρexp.  (g/cm3)  3.74 3.99  ρrel.  (g/cm3)  90.6 98.8  Porosity (vol%)  9.4 1.2  ZrB2 grain size (μm)  2.4 3.6  ρtheor.  (g/cm3)  4.13 4.04  2  \\x0c', 'particles, along with SiC, were observed at the fibre/matrix interface which originated from the reduction of impurity oxides (ZrO2 and SiO2) located in the proximity of the carbon fibres [23,24].  SiO2  density values were adjusted to take into account for the final phases composition of the material which were determined by image analysis. Phases volumetric amounts were normalized on the entire composite (e.g. if the matrix has a 5 % SiC and the matrix constitutes 50 % of the composite, then SiC will only amount to 2.5 % of the entire composite)  3.1.1. Sample ZS Fibres were homogeneously distributed in the ceramic matrix (Fig. 1a), while SiC particles did not show any sign of coarsening, retaining their original particle size < 0.45 μm and seldom forming aggregates < 1 μm (Fig. 1b). From the fibre regions, no evidence was found to indicate a strong chemical reaction between fibre and matrix as the fibres retained their original round shape (Fig. 1c). Some ZrC  \\x0c', 'A. Vinci, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 3. High magnification SEM micrographs of the UHTC matrix of ZSY showing the Y-B-C-O phases with varying stoichiometry and the respective EDS spectra collected at 5 keV. The light grey and black phases are ZrB2 and SiC, the white particles are ZrC and the dark grey phases are Y-B-C-O phases. Si signal was occasionally detected in these phases.  Table 2 4-point flexural strength values from RT to 1500 °C and fracture toughness evaluated with the chevron notch beam test of ZS and ZSY.  Sample  ZS ZSY  σ (MPa)  283 ± 23 436 ± 20  σ1200°C (MPa)  σ1500°C (MPa)  KIc (MPa·m0.5)  - 607 ± 23  - 709 ± 88  8.0 ± 0.9 11.5 ± 0.7  efficacy of carbon is well evident at the fibre/matrix interface where we can recognize a localized area of about 1 μm thick fully densified, while far from the fibre the matrix is more porous. Moreover, we cannot exclude the local formation of small amounts of liquid phase originating from residual silica and boria present on the starting powder particles.  3.1.2. Sample ZSY This composite was characterized by a significantly lower porosity (1.2 %) than ZS (9.4 %) owing to the formation of a lower melting liquid phase that aided the sintering of ZrB2. This was accompanied by an increase of ZrB2 grain size of 50 % (from 2.4 to 3.6 μm) [18]. The effect of Y2O3 as a sintering aid was already reported in the work of. Kovácová et al. [25] where the presence of Y2O3 allowed to increase the relative density from 97.1-98.3. In the present paper the magnitude of Y2O3 addition on the relative  Fig. 5. Load displacement curves for measured at RT, 1200 °C and 1500 °C.  the  4-point flexural  strength of  ZSY  density is much higher. The matrix was nearly fully dense, the ZrB2 grain coarsening suggests dissolution and re-precipitation mechanisms and additional phase formation is observed at the matrix/fibre interface (Fig. 2b and d), in contrast with what reported in the above mentioned paper where no significant change in grain size was observed. All these  Fig. 4. Load displacement curves for the 4-point flexural strength and fracture toughness of ZS and ZSY measured at room temperature.  4  \\x0c', 'A. Vinci, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 6. Fracture surfaces of ZS and ZSY after bending tests at RT: a,d) fracture surfaces, b,e) High magnification of the fibre, c,  f) Fibre/matrix interface.  mechanisms indicate a liquid phase sintering mechanism. According to the SiO2-Y2O3 phase diagram [26], a eutectic liquid phase may form at 1650 °C, between yttria and silica impurities. This liquid was also likely enriched with B2O3 present on the surface of ZrB2 particles. The miscellaneous Y-Si-B-O liquid phase spread in the matrix and at the fibre/matrix interface, helping rearrangement and dissolution of ZrB2, diffusion and re-precipitation. The improvement of densification for ZrB2-SiC systems doped with Y2O3 was reported by other authors and was attributed to the ability of Y2O3 to form liquid phases with the oxide impurities present on the boride and carbide particles, resulting in the strengthening of grain boundaries [18]. However, in the present study, the presence of carbon fibres complicated the picture. Indeed, during re-precipitation from the liquid phase, different phases could form depending on the local chemistry, e.g. the local availability of C or O could cause the formation of prevalently oxides, prevalently carbides or mixed oxy-carbides, as observed by EDS analyses. On closer inspection (Fig. 3), these phases were mainly comprised of Y, B, C and O with varying ratios, while Si signal was hardly found. These phases were characterized by a lamellar structure with features similar to rare earth borocarbides such as YB2C2 (Tm > 2200 K) [27-30] and were likely originated from the reduction of Y2O3 and the B2O3 present on the surface of ZrB2 particles with the carbon of the fibres or fibre debris [29,28]. XRD analysis was carried out on the as produced composite (not shown); the main phases were indexed to ZrB2 and SiC 6H, but minor unidentified phases were observed at high 2θ. However further analysis is needed for a more accurate assessment of these phases structure and characteristics.  3.2. Mechanical properties  The values of strength and fracture toughness at room temperature were 283 MPa and 8.00 MPam°·5 for ZS and 436 MPa and 11.5 MPam0.5 for ZSY respectively (Table 2). The bending strength of fibre-reinforced UHTC composites was found to be lower than the corresponding bulk ceramics [31]. This could be attributed to micro-cracks generated during cooling from the densification temperature due to the CTE mismatch between the fibres and the matrix [32] or the early interlaminar shear failure of the specimens under bending. The strength values obtained in this work were used for comparison purposes. The strength and fracture toughness of ZS were comparable with those of UHTCMCs studied in previous works [33,34] that were in the range of 280-350 MPa and 8-11 MPa m0.5, respectively, whereas the properties of ZSY were significantly higher (436 MPa and 11.5 MPa m0.5 respectively). The higher performance of ZSY can be attributed to the stronger grain boundaries and denser matrix, as well as the slightly higher fibre content. The slope of the load-displacement curves relative to the flexural strength of ZS experiences a small decrease which is typically attributed to a weak fibre/matrix interface and the premature failure of the matrix (Fig. 4). Initially stresses were mostly concentrated in the ceramic matrix. With the increase of the applied load, cracks started to open in the ceramic matrix and the load was transferred to the fibres. For ZSY this effect is negligible, indicating a stronger fibre/ matrix interaction, whereas for ZS some interlaminar failure was observed. ZSY strength was further investigated at high temperature. Strength  5  \\x0c', 'A. Vinci, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 7. X-Ray diffraction patterns of samples ZS and ZSY after oxidation at 1500 and 1650 °C in air. The main peaks are relative to monoclinic ZrO2 (PDF#86-1449), SiO2 (PDF#83-2187). Small peaks relative to YBO3 formation were observed for sample ZSY (PDF#88-0356).  increased from 436 to 607 MPa at 1200 °C; this was attributed to the relaxation of residual stresses accumulated during hot pressing [11-33] and to the increase of the fibre strength at high temperature [35,36]. At 1500 °C, strength increased further to 709 MPa, which is the highest value ever reported for this class of materials. A slight plastic deformation was observed in proximity of the maximum applied load (Fig. 5). The calculated strength at the proportional limit was 516 MPa, which was slightly lower than the strength registered at 1200 °C but still higher than the RT value, while retaining an ultimate strength of 709 MPa. Compared to previously studied composites based on a TaC and HfC matrix, ZrB2 based composites yielded at lower temperatures [11,12]. This could be attributed to the presence of residual low melting phases deriving from the sintering process [18]. In any case, the strength reached is well above the RT value. Looking at the fracture surfaces of ZS and ZSY, limited fibre pull-out was observed in both specimens (Fig. 6a and d), amounting to only 5−20 μm. ZSY displays a more compact and dense ceramic matrix (Fig. 6b and e), but in both composites the first layers of the pitch fibre were anchored to the ceramic matrix (Fig. 6c and f). In the case of ZSY this effect was more dominant and in accordance with the load-displacement curves discussed before.  3.3. Oxidation tests  X-Ray diffraction analysis was carried out after testing at 1500 and 1650 °C in order to identify the species that formed on the surface during oxidation (Fig. 7). In all XRD patterns the highest peak was attributed to hydrated boron oxide, B(OH)3 (PDF#30-0199), which formed on the surface of the specimens after testing due to contact with air humidity. The main phase after oxidation was monoclinic ZrO2 for  both specimens (PDF#86-1449). No tetragonal or cubic ZrO2 were detected. In another work on bulk ZrB2/SiC/Y2O3 it was reported how oxidation carried out at 1500-1650 °C led to the formation of cubic ZrO2 promoted by the presence of Y2O3. However, in the composites studied in our paper, no Y2O3 was left after sintering and only YBC phases were detected. A small peak relative to silicon species (PDF#832187) was detected at 1500 °C for both samples, but it is not visible at 1650 °C, likely due to the partial evaporation and the overall low content of the silica. For ZSY, additional peaks were indexed to YBO3 formation (PDF#88-0356). Following XRD characterization, SEM analysis was carried out on the surface and cross section of the oxidized samples. After testing at 1500 °C, the surface of ZS is characterized by small ZrO2 grains embedded in a glassy silica layer (Fig. 8a), while that of ZSY is characterized by large ZrO2 grains surrounded by a glassy phase containing Y, B, O (Fig. 8d) and small amounts of Si, which was attributed to the formation of YBO3 originating from the oxidation of the Y-B-C-O phases found in the bulk composite. The cross-section of ZS is characterized mainly by two regions as reported previously by the same authors [13,14]: an outer silica layer and an intermediate scale of columnar ZrO2 + SiO2 (Fig. 8b). For ZSY only one scale was observed, mainly consisting of ZrO2 grains held together by a glassy phase of SiO2 and YBO3 (Fig. 8e). This glassy phase was more abundant near the surface. In the case of ZS, the outer fibres were removed due to the oxidation of the carbon to CO (Fig. 8c), while for ZSY the outer fibres were still mostly intact due to the rapid action of the borate that quickly protected them from oxidation (Fig. 8f). The thickness of the oxidized layer of ZS (18 μm) was comparable to that of ZSY (24 μm) and was thinner than that reported in previous works [14]. After oxidation at 1650 °C, the specimens were  visibly more  6  \\x0c', 'and protect the composite from further oxidation, resulting in a slightly lower oxidation resistance than the un-doped specimen. Other works on the kinetics of oxidation of ZrB2/SiC composites showed four main reactions that take place during oxidation [13]:  C + ½ O2  damaged (Fig. 9b and e). The surface of ZS was similar to that observed at 1500 °C but the oxidized layer was thicker (43 μm, Fig. 9c). Columnar zirconia grains grew from the silica melt and were covered with a SiO2 layer (Fig. 9a); this is in agreement with previous works carried out on bulk ZrB2/SiC ceramics [37,38], even though no SiC-depleted layer was observed. This is a phenomenon observed in fibre reinforced UHTC composites where the depleted layer is typically constituted by oxidized fibres [17]. The surface of ZSY was very different: the silica layer, which was barely visible at 1500 °C, was now more abundant on the surface (Fig. 9d), while the intermediate layer consisted mainly of a porous ZrO2 - SiO2 - YBO3 scale (Fig. 9e). Some YBOSi phase was still found at ZrO2 grain boundaries (Fig. 9a,c). The oxidized layer of ZSY increased to 53 μm and was prone to detachment due to an inner porous layer originating from the oxidation of the outer fibres. The morphology of the ZrO2 is also different: for ZS columnar ZrO2 was observed (Fig. 9c), while for ZSY the ZrO2 grains were larger and rounded (Fig. 9f). Previous studies on fibre reinforced ZrB2/SiC with varying SiC contents showed how the oxidation resistance increased with SiC content [14], showing optimal results for SiC > 10 %. Even though Y2O3 played a significant role during the densification process and improved the mechanical properties, it also deprived the composite from SiC which, after sintering, amounted to only 1.1 % of the composite as ascertained by image analysis. The Y-B-O phase was well spread around the ZrO2 grains but there was not enough SiO2 production to fully cover  \\x0c', 'A. Vinci, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 9. Microstructure of ZS and ZSY after oxidation at 1650 °C in air: a,d) surface morphology, b,e) cross-section of the oxidized layer, c,f) detail of the oxidized layer.  which was rather attributed to the overall low SiC content. Very recently, arc-jet tests have shown the capability of these materials to withstand extreme conditions in relevant environments [41].  4. Conclusions  The influence of Y2O3 on the microstructure, thermo-mechanical properties and oxidation resistance of carbon fibre reinforced ZrB2/SiC composites was investigated. Y2O3 was initially added with the goal of improving the oxidation resistance by enhancing the sintering process and promoting the formation of a more compact ZrO2 layer during oxidation. As far as the sintering process is concerned, Y2O3 led to the formation of liquid phases with impurity oxides that aided the sintering of the UHTC phase. However, the presence of carbon in the system led to additional reactions that resulted in the consumption of Y2O3 and formation of yttrium boro-carbides, which was accompanied by the partial removal of SiC. The formation of these novel phases at the fibre/ matrix interface led to very strong interfaces and resulted in higher stiffness and a significant increase of the mechanical properties that were 50 % higher than the undoped composite (strength up to 700 MPa at 1500 °C and fracture toughness of 12 MPam0.5). After oxidation at 1500 °C in air, ZS was characterized by an outer silica layer and an intermediate ZrO2/SiO2 scale, while ZSY was characterized by large ZrO2 grains surrounded by yttrium borate on the surface, and an inner layer of ZrO2/SiO2/YBO3. The oxide layer  the oxi improvement of  thickness was comparable and no significant dation resistance was observed. After oxidation at 1650 °C, the doped sample displayed inferior oxidation resistance; the outer layer was visibly more damaged and the intermediate layer was very porous. This was attributed to the partial evaporation of low melting phases coupled with an overall lower SiC content that resulted in insufficient formation of the protective silica layer. The addition of Y2O3 led to a considerable increase of the relative density and mechanical properties but a decrease of oxidation. Additional tests at higher temperature and in relevant environments, such as those of an Arc-jet wind tunnel are required for the validation of these materials at higher temperatures. A thermo-gravimetric study is currently underway in order to better understand the kinetics of oxidation, while TEM analysis is being carried out in order to better study the structure and role of the YBC phases found.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgements  This work has received funding from the European Union’s Horizon  8  \\x0c', 'A. Vinci, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  2020 “Research and innovation programme” under grant agreement No. 685594 (C3HARME). The authors are grateful to Cesare Melandri for mechanical testing, Daniele Dalle Fabbriche for hot pressing and Mauro Mazzocchi for XRD analysis. The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.  References  [7]  [6]  [4]  [3]  [1]  [10]  S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra-high temperature ceramics for aeropropulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767, https://doi.org/10.1016/S0955-2219(02)00140-1. [2] M.M. Opeka, I.G. Talmy, E.J. Wuchina, Ja. Zaykoski, S.J. Causey, Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. 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  "_id": 108,
  "PDF": "Inhomogeneous oxidation of ZrB 2 -SiC ultra-high-temperature ceramic particulate composites and its mitigation.pdf",
  "Text": "['Acta Materialia 129 (2017) 138e148  Contents lists available at ScienceDirect  Acta Materialia  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / a c t a m a t  Full  length article  Inhomogeneous oxidation of ZrB2-SiC ultra-high-temperature ceramic particulate composites and its mitigation  Lin Zhang, Nitin P. Padture*  School of Engineering, Brown University, Providence, RI 02912, USA  a r t i c l e  i n f o  a b s t r a c t  The ubiquitous inhomogeneous-oxidation behavior of a prototypical ZrB2-20 vol% SiC ultra-hightemperature ceramic (UHTC) particulate composite is investigated in detail. An innovative approach for enhancing the oxidation resistance of the composite and delaying the onset of its inhomogeneous oxidation is demonstrated. In this approach, a borosilicate glass coating that is highly compatible with the composite is applied to the composite surface a priori. The effect of the borosilicate-glass coating thickness on the homogeneous-oxidation kinetics of the composite is investigated quantitatively and modeled. A new phenomenon is discovered where ZrO2 precipitates of graded size are found to nucleate and grow within the borosilicate-glass layer, both in natively grown and in applied coating, during the oxidation of the composites. This is attributed to the ZrO2-solubility gradient across the borosilicate-glass layer as a result of preferential B2O3 volatilization at the surface. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Article history:  Received 15 January 2017 Received in revised form 22 February 2017 Accepted 25 February 2017 Available online 27 February 2017  Keywords:  Ultra-high-temperature ceramics Composites Oxidation Glass Zirconium diboride Silicon carbide  1.  Introduction  Ultra-high-temperature ceramics (UHTCs), which include transition-metal borides, carbide, and nitrides, have potential uses in aerospace and other applications [1e3]. UHTCs have extremely high melting points, which allow them to withstand extreme environments that may exist in applications such as sharp leading edges of hypersonic vehicles [1e4]. These environments are typically oxidizing, making oxidation properties of UHTCs very important [5,6]. In this context, unlike metal-carbide and metalnitride UHTCs, the oxidation of metal-boride UHTCs results in the formation of non-gaseous B2O3 glass as one of the oxidation products, which can protect the UHTCs [5e7]. Among the metalboride family of UHTCs, ZrB2 is the most widely studied as it possesses perhaps the best combination of properties [1,7,8]: high melting temperature, moderate density, high thermal conductivity, and low thermal expansion coefficient. However, ZrB2 UHTCs are brittle [8], and the B2O3 oxidation product is volatile [5,6]. To address these issues, bulk ZrB2 UHTCs are typically reinforced by particulate SiC (20e30 vol%) [5e7,9,10]. This enhances the toughness of UHTC composites, and the oxidation product is a more  * Corresponding author. E-mail address: nitin_padture@brown.edu (N.P. Padture).  http://dx.doi.org/10.1016/j.actamat.2017.02.076 1359-6454/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  viscous, less volatile borosilicate glass [5,6]. The oxidation of ZrB2SiC UHTC composites typically results in a multi-layer structure. This includes an outermost borosilicate-glass layer, with an oxide layer underneath comprising a porous ZrO2 skeleton filled with the borosilicate glass [5,6,11]. In some cases, the existence of another layer d ZrB2 depleted in SiC d between the oxide layer and the base ZrB2-SiC UHTC composite has been reported [11e13]. One of the challenges with the oxidation of ZrB2-SiC UHTC composites is the extreme inhomogeneity of the multi-layer structure. Large variability in the thicknesses of the different layers and in the weight changes is routinely observed in identical specimens/materials subjected to identical oxidation conditions [14,15]. For example, Shugart et al. [15] have shown up to 80% variation in the average thickness of the oxide layer in identical ZrB2-SiC UHTC composites. This inhomogeneous oxidation behavior has been attributed to the prevalence of stochastic chemical and physical processes, which invariably involve the borosilicate-glass layer [14e18]. Here we have investigated the inhomogeneous oxidation behavior of a fully-dense ZrB2-20 vol% SiC UHTC composite, and confirm the critical role played by the borosilicate-glass layer in this behavior. We have also discovered that the borosilicate-glass layer contains small ZrO2 precipitates of graded size, which appear to nucleate and coarsen within the glass. More importantly, we show that the application of a pre-determined coating of the borosilicate  \\x0c', \"L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  139  glass on the ZrB2-SiC UHTC composite surface before oxidation enhances significantly its oxidation resistance, and helps mitigate the inhomogeneous oxidation behavior. A theoretical model is built, based on the previous work of Parathasarathy et al. [19], to understand and to predict the effect of the borosilicate-glass layer thickness on the oxidation behavior of ZrB2-SiC UHTC composites.  2.  Experimental procedure  2.1.  ZrB2-SiC UHTC composites fabrication  Commercially available starting powders were used: ZrB2 powder (average particle size ~2 mm, Grade B, H.C. Starck Corp., Newton, MA) and SiC (average particle size ~1.7 mm, Grade B-hp, bSiC, H.C. Starck Corp., Newton, MA). Batches composed of 80 vol% ZrB2 powder and 20 vol% SiC d the most popular composition d were ball-milled, using zirconia-balls media in ethanol, for 24 h. The resulting slurries were dried, while being stirred, to obtain homogenously dispersed powder mixtures, which were then crushed in a mortar-pestle and sieved. Small batches of the powder mixture (3e4 g) were place in graphite dies (20 mm diameter) lined with graphite foil, and densified using spark plasma sintering (SPS; Thermal Technologies LLC, Santa Rosa, CA). The SPS conditions were: 1850 \\x0eC peak temperature, 75 MPa pressure, 5 min hold duration, and 100 \\x0eC.min heating/cooling ramp rate. The resulting specimens were ground and cleaned. Their densities were measured using the Archimedes' method, with distilled water as the immersion medium. The theoretical density of the \\x003 based on rule of mixtures, using density composite is 5.51 Mg m \\x003 for ZrB2 and SiC, respectively. values of 6.085 and 3.210 Mg m These disk specimens were further machined for the various investigations. Cross-sections were cut from the disks and polished to a 1-mm finish using routine metallographic techniques for observing the microstructures in the scanning electron microscope (SEM; LEO 1530VP, Carl Zeiss, Munich, Germany). Rectangular parallelepipeds of size 4 \\x02 3 \\x02 2 mm3 were cut from the discs for oxidation studies of the as-sintered ZrB2-SiC UHTC composites. Oxidation studies with the pre-determined coating of the borosilicate glass applied on the surface of ZrB2-SiC UHTC composites were performed on discs that were cut in half (semicircular).  \\x001  2.2.  Borosilicate glass processing  Borosilicate glass for application on the surface of ZrB2-SiC UHTC composite as a coating was prepared separately. Commercially available B2O3 (~40 mesh, Acros, Geel, Belgium) and SiO2 (particle size of 1e5 mm, Atlantic Equipment Engineers, Bergenfield, NJ) powder were mixed in the molar ratio of 3:2, and ball-milled for 24 h (zirconia-balls media in ethanol), and then dried on a hot-plate while stirring. This composition is typical of the borosilicate glass that forms after the oxidation of ZrB2-SiC UHTC composite [19,20]. The powder mixture was melted at 1000 \\x0eC for 2 h in a platinum crucible in air using a box furnace (Thermolyne, Dubuque, IA) and then quenched to room temperature in water. The final borosilicate glass powder was formed by crushing the glass and ball-milling it for 24 h (zirconia-balls media in ethanol), and then dried on hotplate. The weight loss after melting was found to be negligible, confirming that the B2O3:SiO2:3:2 molar ratio is unchanged. The amorphous nature of this borosilicate glass powder was confirmed using X-ray diffractometry (XRD; D8 Advance, Bruker AXS, Karlsruhe, Germany), where only a glassy hump is observed (not shown here).  2.3. Oxidation testing  The rectangular-parallelepiped specimens of the as-sintered ZrB2-SiC UHTC composite were placed at ~45\\x0e angle in a zirconia crucible against its wall such that only two edges of the specimens were in contact with the crucible (floor and wall). These specimens were oxidized at 1500 \\x0e C for 10 min to 5 h in air in the box furnace. The specimens were weighed before and after each oxidation heattreatment. Two cross-sections, one near the top and one near the bottom, were cut from each of the oxidized rectangularparallelepiped specimens. They were polished to a 1-mm finish and observed in an optical microscope (Eclipse LV100, Nikon, Tokyo, Japan) and the SEM. The thicknesses of the different oxidation layers were measured at various locations using image analysis in both cross-sections. This approach allowed us to investigate quantitatively the inhomogeneous oxidation behavior of all four faces of the specimens, and also any possible effect of gravity. To study the oxidation behavior of the ZrB2-SiC UHTC composite with the applied coating of the borosilicate glass, semicircular specimens (20 mm diameter) were used. Here the specimen was laid flat on a zirconia plate, and the borosilicate glass powder in the form of a paste (mixed in ethanol) was applied over a small circular area at the center of the top surface. The amount of the borosilicate glass was controlled so as to result in ~50 mm, ~100 mm, or ~200 mm thick coating in the center of the circular area after melting. (The density of the borosilicate glass is assumed to be 1.95 Mg m Ref. [21].) Three specimens for each borosilicate-glass coating thickness were prepared. The oxidation tests were performed at 1500 \\x0e C in air, but for various durations (0.5e20 h) for each of the specimens. Control bare specimens without the applied borosilicate-glass coating were also subjected to oxidation heattreatments. All the oxidized specimens were cut in the middle, and the cross-section were polished to a 1-mm finish and observed in the SEM. The thickness of the oxide layer at and near the center of the borosilicate-glass circular coating was measured on the SEM micrographs; 10 measurements per specimen. Compositional (elemental) maps of the oxidation layers were obtained in the SEM using energy dispersive spectroscopy (EDS). A transmission electron microscope (TEM; 2100F, JEOL, Peabody, MA), operated at 200 keV accelerating voltage and also equipped with EDS and electron energy loss spectrometer (EELS), was used to characterize the cross-sections of some samples. TEM specimens were extracted from specific locations by in situ lift-out of foils nano-machined using a focused ion beam microscope (FIB; Helios 600, FEI, Hillsboro, OR). Selected area electron diffraction patterns (SAEDPs) were collected from various grains and indexed using standard procedures.  \\x003  3. Results  Fig. 1 shows a cross-sectional SEM image of a polished crosssection of the ZrB2-SiC UHTC composite. The full density of the material and the uniform dispersion of the SiC particulates are evident in this micrograph. The density was measured to be \\x003 or 99.5% of the theoretical 5.48 Mg m limit. Fig. 2A shows the specimen geometry and its placement in the crucible for the oxidation tests. Fig. 2B is an example of a crosssectional optical image of the oxidized ZrB2-SiC UHTC composite (1500 \\x0e C, 5 h, in air), showing highly inhomogeneous oxidation behavior. Fig. 2C presents quantitative data showing large variability in the thicknesses of the borosilicate-glass and the oxide layers. No correlation between any of these data could be found, indicating random variability regardless of the side of the rectangular-parallelepiped and no effect of gravity. Fig. 3 presents  \\x0c\", '140  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  Fig. 1. Cross-sectional SEM image of ZrB2-SiC UHTC composite. The grey phase is ZrB2 and the dark phase is SiC.  Fig. 3. Weight gain of ZrB2-SiC UHTC composite as a function of oxidation duration at 1500 \\x0e C in air.  Fig. 2. (A) Schematic diagram showing the specimen geometry and its placement in the crucible. Cross-sections 1 and 2 are indicated. (B) Optical image of polished cross-section 1 of ZrB2-SiC UHTC composite specimen oxidized at 1500 \\x0e C for 5 h in air. (C) Thicknesses of the borosilicate-glass layer and the oxide layer for the four surfaces of the oxidized specimen at cross-sections 1 and 2 (1500 \\x0e C, 5 h, in air). Each gray bar represents an average of 10 measurements along the length of a side at regular intervals, and the error bar represent one standard deviation.  the weight-gain data, showing a near-parabolic time dependence. While these data are included for the sake of completeness, not much can be inferred from them because of the complex oxidation  behavior in ZrB2-SiC UHTC composites. Fig. 4A and B shows photographs of semicircular disk specimen the ZrB2-SiC UHTC composite with circular borosilicate-glass  of  \\x0c', 'L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  141  Fig. 4. ZrB2-SiC UHTC composite coated with ~100 mm borosilicate glass in a circular area in the center: (A) before and (B) after oxidation at 1500 \\x0e C for 1 h in air.  coating (~100 mm) applied in the center, before and after the (1500 \\x0eC, 1 h, air), oxidation heat-treatment respectively. These photographs confirm that circular shape of the borosilicate-glass layer is maintained, with negligible flow of the glass. The white specks outside of the circle in Fig. 4B are possibly remnants of the random ‘convection cells,’ which are discussed later [14]. No such specks are visible in the center of the circular region. Fig. 5 compares the oxidation behavior (1500 \\x0eC, 4 h, in air) of the ZrB2-SiC UHTC composite without and with the applied borosilicate-glass coating (~50 mm), showing clearly a thick, heterogeneous oxide layer (25e50 mm) in the former (Fig. 5AeD) and a thin, homogenous oxide layer (~10 mm) in the latter (Fig. 5EeH). The higher magnification SEM images in Fig. 6A and B highlight this difference in the oxidation behavior even more clearly. With thicker applied borosilicate-glass coating (~100 mm and ~200 mm), the uniformity in the oxide layer over a large area is remarkable, as shown in the low-magnification SEM images in Fig. 6C and D. These results show that the application of the borosilicate-glass coating increases the oxidation resistance of the ZrB2-SiC UHTC composite, and mitigates the inhomogeneous oxidation behavior. A look at the region near the edge of the tapering borosilicate-glass layer in Fig. 7A and B further confirms this conclusion. As the borosilicateglass coating becomes vanishingly thin near the edge, the oxide layer goes from being uniform (on the right) to highly inhomogeneous (on the left). Note that the oxide layer is similar to what has been observed by others [5,6]. d a porous ZrO2 skeleton filled with the borosilicate glass. The SiC-depleted ZrB2-SiC layer was not observed in this case. The oxide-layer thicknesses (homogenous oxidation behavior) as a function of oxidation duration (1500 \\x0eC, in air) for the three different applied borosilicate-glass coating thicknesses are plotted in Fig. 8AeC. The improved oxidation resistance is clearly evident from these data. Note that oxide-layer thickness data for longer oxidation durations are not plotted in Fig. 8AeC because even with an applied borosilicate-glass coating eventually inhomogeneous oxidation behavior ensues, making the oxide-layer thickness data less meaningful. For example, the cross-sectional SEM image in Fig. 9 shows inhomogeneous oxidation at 20 h (1500 \\x0eC, in air) in ZrB2-SiC UHTC composite even with a ~100-mm borosilicate-glass applied coating. The solid curves in Fig. 8AeC are predications from an oxidation model described in the next section. The borosilicate-glass layer (both native and with applied coating) was examined in more detail using the TEM, which has led to the discovery of an interesting phenomenon. Fig. 10A is a crosssectional bright-field TEM image of the region between the highly inhomogeneous oxide layer and the borosilicate-glass layer in an oxidized (1500 \\x0e C, 0.5 h, in air) ZrB2-SiC UHTC composite without the applied borosilicate-glass coating. Detailed TEM studies in Fig. 10D reveal that very close to the oxide layer there is clean borosilicate glass, but just above that region one observes very small precipitates (~50 nm). In the middle of the borosilicate-glass layer (Fig. 10C), the size of the precipitates becomes larger (~100 nm).  These particles contain Zr and O, as confirmed using EDS (Fig. 10F) and EELS (Fig. 10G), respectively. Isolated larger precipitates (~300 nm in size) are found only near the top, and are confirmed to be m-ZrO2 using SAEDP (Fig. 10E). The same phenomenon is also observed in the borosilicate-glass layer above homogenous oxide layer (Fig. 11AeC) where convection cells are absent. Furthermore, this phenomenon is observed in ZrB2-SiC UHTC composite coated with borosilicate glass (~100 mm thickness) and oxidized at 1500 \\x0eC for 1 h in air (Fig. 12).  4. Oxidation model  To understand better, and to predict, the oxidation kinetics of ZrB2-SiC UHTC composites with applied borosilicate-glass coating, a model is developed here, one that builds on the previous work of Parathasarathy et al. [19] Fig. 13 shows a schematic diagram of the idealized oxidation behavior of ZrB2-SiC UHTC composites used in for intermediate temperatures (<1700 \\x0e C) where SiO2 this model volatilization is negligible. As mentioned earlier, the oxidation of ZrB2-SiC UHTC composites results in several layers, which is confirmed experimentally: the outermost glass layer (l3a ), the intermediate oxide layer (l23 ), and the substrate (l12 ). At the interface of the substrate layer and the oxidized layer (denoted by i2), the oxidation occurs through two main reactions as follows:  ZrB2 ðsÞ þ 5 2  O2 ðgÞ ¼ ZrO2 ðsÞ þ B2O3 ðlÞ  SiC ðsÞ þ 3 2  O2 ðgÞ ¼ SiO2 ðlÞ þ COðgÞ  (1)  (2)  Explicit assumption in this model is that the oxygen transport to the interface i2 (composite surface) controls the oxidation kinetics [19]. The partial pressure of oxygen at the interface i2 can be calculated using the equilibrium constant for ZrB2 oxidation [19].  (\\x00 \\x00  \\x01\\x00 \\x01\\x00  \\x01 )2 \\x01  5  PO2 \\x00i2 ¼  aZrO2 aZrB2  aB2 O3 KZrB2  ;  (3)  where PO2\\x00i2 is the oxygen partial pressure at i2, KZrB2 is the equilibrium constant, and aZrO2 , aB2 O3 , and aZrB2 are the activities of ZrO2, B2O3 and ZrB2, respectively. Assuming the activity of the solid phases are equal to unity, the activity of B2O3 at interface i2 can be obtained in terms of the molar ratio of ZrB2 and SiC [19].  ¼  aB2 O3  (4)  \\x11  \\x10 \\x11  ð1 \\x00 fs Þ  \\x10  rZrB2  ðMZrB2  Þ  ;  ð1 \\x00 fs Þ  rZrB2  ðMZrB2  Þ þ fs  ðrSiC Þ  ðMSiC Þ  where fs  is the volume fraction of SiC in the composite; rZrB2 and  \\x0c', '142  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  Fig. 5. Cross-sectional SEM images and corresponding EDS elemental maps of ZrB2-SiC UHTC composite oxidized at 1500 \\x0e C for 4 h in air: (AeD) without and (EeG) with applied borosilicate-glass coating (~50 mm).  rSiC are the densities of ZrB2 and SiC, respectively; and MZrB2 and MSiC are the molecular weights of ZrB2 and SiC, respectively. Since no SiC-depleted layer is observed experimentally, it is assumed that there is no preferred oxidation of ZrB2 or SiC. Thus, the fraction of oxygen reacting with ZrB2 and SiC at i2, (fO2\\x00ZrB2\\x00i2 ) and (fO2\\x00SiC\\x00i2 )  fO2\\x00ZrB2\\x00i2 ¼  5 2  \\x00  5 2  fZrB2  respectively, can be calculated as follows:  \\x01  \\x00  \\x01 þ 3  fZrB2  2  ðfSiC Þ  (5)  \\x0c', 'L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  143  Fig. 6. Cross-sectional high-magnification SEM images of ZrB2-SiC UHTC composite oxidized at 1500 \\x0e C for 4 h in air: (A) without and (B) with applied borosilicate-glass coating (~50 mm). Cross-sectional low-magnification SEM images of ZrB2-SiC UHTC composite oxidized at 1500 \\x0e C in air with applied borosilicate-glass coating: (C) ~100 mm (7 h) and (D) ~200 mm (4 h).  the oxygen fluxes through the oxide layer (JO2 ð23Þ ) and the glass layer (JO2 ð3aÞ ) are equal. They are related to the respective oxygen partial pressure gradients and oxygen permeability coefficients (p) of the borosilicate glass, and are given by [19].  (PO2 )  \\x00  \\x0c\\x0c\\x0cJO2 ð23Þ  \\x0c\\x0c\\x0c ¼  \\x01  \\x10  \\x01 \\x00 \\x00  l23  PO2\\x00i3  PO2\\x00i2  pO2\\x00glass\\x0023  1 \\x00 fZrO2  \\x11\\x00  \\x01  and  \\x0c\\x0c\\x0cJO2 ð3aÞ  \\x0c\\x0c\\x0c ¼  \\x00  PO2\\x00a  giving:  \\x01 \\x00 \\x00  l3a  PO2\\x00i3  \\x01  \\x10  \\x11  pO2\\x00glass\\x003a  ;  (8)  (9)  (10)  fO2\\x00SiC\\x00i2 ¼ 1 \\x00 \\x00  \\x01  fO2 \\x00ZrB2\\x00i2  ;  (6)  where fZrB2 and fSiC are the mole fractions of ZrB2 and SiC, respectively, in the composite. The rate of growth of the oxide-layer thickness (l23 ), which identifies with the ZrO2 skeleton size, is then given by [19].  \\x1a  \\x10  \\x11\\x00  dl23  dt  ¼ VZrO2  2  5  JO2 ð23Þ  fO2\\x00ZrB2\\x00i2  \\x01\\x1b  \\x01 ;  \\x00  1  fZrO2  (7)  where VZrO2 is the molar volume of ZrO2, fZrO2 is the fraction of ZrO2 in the oxide layer, and JO2 ð23Þ is the flux of oxygen through the oxide layer. Since the only path for oxygen to reach the composite is to diffuse through the liquid borosilicate glass, the absolute values of  n  \\x10  l23  PO2\\x00i3 ¼  \\x11\\x00  n  \\x10  pO2\\x00glass\\x003a  PO2\\x00a  l23  pO2\\x00glass\\x003a  \\x01o \\x11o  nh n\\x10  fg l3a  \\x10 \\x10  pO2\\x00glass\\x0023  \\x01io  \\x11\\x00 \\x11\\x11o  PO2\\x00i2  fg l3a  pO2 \\x00glass\\x0023  :  þ  þ  \\x0c', '144  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  Fig. 7. Cross-sectional SEM images of ZrB2-SiC UHTC composite near the tapered edge of the applied borosilicate-glass coating (~100 mm in the center) oxidized at 1500 \\x0e C in air for: (A) 1 h and (B) 4 h. The dashed curve indicates the interface between the oxide layer and the unaffected composite.  The oxygen permeability coefficients depend strongly on the composition of the borosilicate glass, and they were calculated by log-mean interpolation of the values for pure SiO2 and B2O3 end points [19].  pO2\\x00B2 O3  and pO2\\x00SiO2  \\x12\\x0016000 ¼ 0:15 exp \\x12\\x0020808 \\x007 exp  ¼ 10  T  \\x13 \\x13  T  :  (11)  pO2 \\x00glass\\x003a  Since the outermost borosilicate glass layer (l3a ) is silicon-rich due to the preferential volatilization of B2O3, compared to the borosilicate glass in the oxide layer, differs from pO2\\x00glass\\x0023 . Assuming B2O3 volatilization does not occur at the interior i2 interface, B2O3 concentration (mole fraction) in the borosilicate glass at that interface is 0.72 for 20 vol% SiC in the composite [19]. The B2O3 concentration in the outermost glass layer is assumed to be equal to that in the applied borosilicate glass layer (0.6), which is close to what has been used by others [19,20]. The rate of change of the thickness of the outermost borosilicate glass layer (l3a ) depends on the rate of volume increase of liquid B2O3 and SiO2, and the rate of consumption of the glass in the oxide layer (l23 ), where the former depends on the rate of evaporation losses (JB2 O3 and JSiO2 ) [19].  Fig. 8. Oxide-layer thickness as a function of oxidation time in ZrB2-SiC UHTC composite oxidized at 1500 \\x0e C in air with applied borosilicate-glass coating: (A) ~50 mm, (B) ~100 mm, and (C) ~200 mm. The solid squares are experimental data points (three specimens per condition, 10 measurements per specimen), and the error bars represent one standard deviation. The solid curves are predictions of the oxidation model.  \\x13 \\x12  \\x13  ¼  dl3a  dt  dl23 dt  fZrO2 \\x01 \\x00 JSiO2 VZrO2  \\x00 JB2 O3  \\x1b  VSiO2  JO2 ð23Þ  \\x1a\\x12  \\x02 \\x00  \\x1b  VB2 O3  \\x00  \\x00 dl23  dt  \\x1a  2  \\x00  3  þ  1 \\x00 fZrO2  \\x01  fO2\\x00SiC\\x00i2  \\x01  :  (12)  The evaporation rate of B2O3 and SiO2 can be obtained using [19].  \\x0c', 'L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  145  Fig. 9. Cross-sectional  low-magnification SEM image of ZrB2-SiC UHTC composite with ~100 mm borosilicate-glass applied coating oxidized for 20 h at 1500 \\x0e C in air.  Fig. 10. Cross-sectional bright-field TEM images of ZrB2-SiC UHTC composite oxidized at 1500 \\x0e C for 0.5 h in air: (A) low-magnification image of the inhomogeneous oxide layer and borosilicate-glass layer (FIBed TEM specimen); (B)e(D) high-magnification images as indicated; (E) indexed SAEDP from the m-ZrO2 particle in (B), with transmitted beam (T) and zone axis (B) indicated; (F) EDS elemental Zr-map of (C); and (G) EELS spectrum from a ZrO2 particle in (C) showing the peak associated with O. The Pt coating in (A) is from the FIBing process.  Jspecies ¼ Dspecies Pspecies  RT dB  ;  !1  2  !1 6 \\x00  where dB is the boundary layer thickness given by:  dB ¼ 3 2  lC  nfluid  hfluid rfluid  \\x011  Dspecies  3 ;  (13)  (14)  with Dspecies and Pspecies being the diffusion coefficient and the vapor pressure of the different species, respectively; is the composite thickness; nfluid , hfluid and rfluid are the velocity, the viscosity, and the density of the fluid, respectively; and R and T have their usual meanings. The thicknesses of the different layers as a function of oxidation duration (t) can now be calculated by using Eqs. (3)e(14), and are  lC  plotted in Fig. 8AeC. An infinitesimal number is chosen as the initial value for l23, whereas the initial value for l3a equals to the thickness of the applied borosilicate-glass coating (50 mm, 100 mm, or 200 mm). A list of parameters and their values used in these calculations are those from Parthasarthy et al. [19], and are listed in 50-mm applied Table 1. The trend observed in Fig. 8A for a borosilicate-glass coating is consistent with the experimental data, and predicts parabolic oxidation behavior for longer oxidation durations. The enhanced oxidation-resistance with thicker applied borosilicate-glass coating is clearly captured by the oxidation model in Fig. 8B and C, demonstrating the effectiveness of this approach.  5. Discussion  There are two main mechanisms responsible for the inhomo\\x0eC) geneous oxidation of ZrB2-SiC UHTC composites (at 1500      \\x0c', '146  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  Fig. 11. Cross-sectional high-magnification bright-field TEM images of different regions within the borosilicate-glass layer above homogeneous oxide layer in an oxidized ZrB2-SiC UHTC composite (1500 \\x0e C, 0.5 h, in air): (A) top, (B) middle, and (C) bottom. (D) Schematic illustration of the observed phenomenon in Figs. 10 and 11.  Fig. 12. Cross-sectional high-magnification SEM image of borosilicate-glass layer above homogeneous oxide layer in an oxidized (1500 \\x0e C, 1 h, in air) ZrB2-SiC UHTC composite with an applied borosilicate-glass coating (~100 mm).  Fig. 13. Schematic illustration of the simplified oxidation behavior of ZrB2-SiC UHTC composites used in the oxidation model. Phases, layers, and interfaces are marked.  \\x0c', 'Table 1 Summary of the parameters used in the model [19].  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  147  Symbol  aZrB2 aZrO2 aB2 O3 DB2 O3 DSiO2 rSiC rZrB2 MSiC MZrB2 PB2 O3 PSiO2 VB2 O3 VSiO2 VZrO2 fZrB2 fSiC fO2 \\x00ZrB2 \\x00i2 fO2 \\x00SiC\\x00i2  fS  fZrO2  JO2 ð23Þ JO2 ð3aÞ  JB2 O3 JSiO2 KZrB2  lC  l23 l3a  PO2 \\x00a  PO2 \\x00i3 PO2 \\x00i2  dB  nfluid hfluid rfluid pO2 \\x00glass\\x0023 pO2 \\x00glass\\x003a  Units  e  \\x001  m2 s  \\x003  Mg m  \\x001  kg mol  Atm.  m3 mol  \\x001  e  e  e  mol m  \\x002 s  mol m  \\x002 s  e  m  Atm.  m  m s  \\x001  Pa.s  Mg m  mol m Atm  \\x001  \\x003 \\x001 s  Description  Activity  Diffusion coefficient  Density  Molecular weight  Vapor pressure  Molar volume  Mole fraction  Fraction of oxygen reacting at interface i2  Volume fraction of SiC in the composite  Volume fraction of ZrO2 in the oxide layer Oxygen flux through oxide layer  Oxygen flux through glass layer  B2O3 vapor flux at the outermost surface SiO2 vapor flux at the outermost surface Equilibrium constant  Composite thickness  Oxide-layer thickness  Glass-layer thickness  Oxygen partial pressure at interface a  Oxygen partial pressure at interface i3  Oxygen partial pressure at interface i2  Surface boundary-layer thickness  Fluid velocity  Fluid viscosity  Fluid density  Glass oxygen permeability in oxide layer  Oxygen permeability of glass layer  Value or source  1  1  Eq. (4) 1.18 \\x02 10 1.29 \\x02 10  3.21 4.01 \\x02 10 6.08 1.13 \\x02 10 1.93 \\x02 10 2.83 \\x02 10 3.16 \\x02 10 3.26 \\x02 10 2.17 \\x02 10 0.73  \\x004 \\x004  \\x002 \\x001 \\x003 \\x009 \\x005 \\x005 \\x005  0.27  Eq. (5)  Eq. (6)  0.20  0.77  Eq. (8)  Eq. (9)  Eq. (13)  2.95 \\x02 1034 1.5 \\x02 10 Eq. (7)  \\x003  Eq. (12)  0.21  Eq. (10)  Eq. (3)  Eq. (14)  0.01 1.83 \\x02 10  \\x005  1.23  Eq. (11)  \\x001  reported in the literature [14,15]. The first occurs in the very initial stages (within seconds), where a very thin borosilicate-glass layer forms on the surface due the oxidation of ZrB2 and SiC according to Eqs. (1) and (2), respectively [15]. The gaseous CO, and possibly volatilized B2O3 gas, build up under the very thin borosilicate-glass layer creating gas bubbles that are randomly distributed [15]. It appears that within 30 s, the bubbles burst, exposing fresh ZrB2-SiC UHTC composite surface to oxidation [15]. However, with continued oxidation, and concomitant thickening of the borosilicate-glass layer and the oxide layer, the bursting bubbles do not expose fresh ZrB2-SiC UHTC composite. This may suppress inhomogeneous oxidation for some time. However, as oxidation progresses, it has been shown that the large volume change associated with ZrB2 and SiC oxidation (Eqs. (1) and (2)) results in a convective upward flow of B2O3 -SiO2-ZrO2 through the borosilicate-glass layer and volatilization of B2O3 at the surface [14]. This in turn results in randomly distributed convection cells of fastoxygen paths within the borosilicate-glass layer and the disruption of the oxide layer [14]. The latter also results in the release of ZrO2 particles (several microns) from the oxide layer into the borosilicate glass [14]. Thus, the formation of the convection cells appears to be responsible for the significant inhomogeneity observed in the oxide-layer thickness in Figs. 2, 5AeD, and 6A. This led us to invent the approach of applying a pre-determined coating of borosilicate glass on the ZrB2-SiC UHTC composite surface. The thick borosilicate-glass coating prevents the initial gas-bubble bursting, avoiding inhomogeneous oxidation. More importantly,  beyond the initial stage, the overall oxidation rate and the onset of the formation of convection-cells is delayed significantly with increasing thickness of the applied borosilicate-glass coating. Of course, the use of coatings to improve oxidation resistance of ceramics is not new. However, the possible incompatibility between the coating material and the substrate, and also the oxidation products, is an overriding issue that always needs to be addressed. In this context, our approach uses the oxidation-product material itself as the coating for better compatibility d the simplicity and the effectiveness of this approach are clearly demonstrated. While borosilicate glass of a specific composition was applied using a simple paste-application method here, other compositions and application methods could be even more effective. The phenomenon presented in Figs. 10e12 is scientifically interesting, and it has not been investigated in detail before as there is a paucity of TEM studies of the borosilicate-glass layer in the literature. (Note that Karlsdottir et al. [14] have reported a crosssectional SEM image showing layers fine particles within the borosilicate-glass layer.) The gradient in the size of the ZrO2 precipitates, increasing from the oxide-layer to the outermost surface, is unmistakable. Also, the existence of the ZrO2-precipitates-free zone closest to the oxide layer is consistent in Figs. 10D, 11C and 12. The ZrO2 precipitates are much too small to have come from the convection cells, where in the latter the ZrO2 particles are at least an order of magnitude larger [14]. They are not likely to have come from the bursting of the bubbles in the initial stages of oxidation either because this phenomenon is also observed in the case of  \\x0c', '148  L. Zhang, N.P. Padture / Acta Materialia 129 (2017) 138e148  ZrB2-SiC UHTC composite with a pre-deposited ~100 mm borosilicate-glass coating (Fig. 12). The most likely mechanism for the origin of this phenomenon is as follows. The preferential volatilization of B2O3 at the outermost surface of the borosilicateglass layer sets up a B2O3-concentration gradient across the layer. Shugart et al. [20] have indicated that the B2O3 concentration decreases from ~60 mol% in the oxide layer to a negligible value at the outermost surface of the borosilicate-glass layer. It is known that ZrO2 has moderate solubility in B2O3-rich borosilicate liquid at 1500 \\x0eC, as shown in Fig. 14 [16]. However, as the B2O3/SiO2 ratio in the liquid decreases, the ZrO2 solubility decreases substantially, going to near zero in a SiO2-rich borosilicate liquid. Thus, in the zone just above the oxide layer (see e.g. Fig. 11D) the solubility of ZrO2 in the borosilicate glass is expected to be lower due to some B2O3 depletion. However, any supersaturated ZrO2 is likely to deposit on the ZrO2 grains in the oxide layer which are in such close proximity, resulting in the ZrO2-precipitate-free zone. As one moves up the borosilicate-glass layer, the B2O3/SiO2 ratio and the ZrO2 solubility are lower, but now the ZrO2 grains in the oxide layer are relatively further away. Thus, the supersaturated ZrO2 in that region is likely to nucleate homogenously as very small ZrO2 precipitates. As one moves further up, the ZrO2 supersaturation increases due to further B2O3-depletion, resulting in the growth of the ZrO2 precipitates. Since this phenomenon depends only on the B2O3-concentration gradient in the borosilicate-glass layer and the presence of ZrO2 grains in the oxide layer, it is observed regardless of whether the oxide layer is inhomogeneous (Fig. 10) or homogeneous (Fig. 11), or whether the borosilicate-glass layer forms natively (Figs. 10 and 11) or it is deposited as a coating before oxidation (Fig. 12).  6.  Summary  Detailed oxidation (1500 \\x0e C, in air) studies of the fully-dense ZrB2-20 vol% SiC UHTC particulate composite confirm quantitatively its inhomogeneous oxidation behavior. The outermost borosilicate-glass layer that forms as a result of the oxidation is closely involved in the inhomogeneous oxidation of the ZrB2-SiC UHTC composite. The application of a pre-determined coating of a  Fig. 14. Calculated 1500 \\x0e C section of the SiO2-B2O3-ZrO2 ternary phase diagram showing the extent of the liquid region. Reprinted with permission from Ref. [16]. Arrows indicate the reduction of ZrO2 solubility in the liquid with B2O3 depletion and SiO2 enrichment.  borosilicate glass on the surface of the UHTC composite before oxidation is found to improve the oxidation resistance profoundly, and it is found to mitigate the inhomogeneous oxidation behavior by delaying significantly its onset. Results from the theoretical model built upon the previous work of Parathasarathy et al. [19], captures the effect of the borosilicate-glass layer thickness on the oxidation behavior of the ZrB2-SiC UHTC composite. The borosilicate-glass layer in all cases is found to contain small ZrO2 precipitates of graded size, which appear to nucleate homogeneously above the oxide layer and coarsen as one moves toward the surface of the borosilicate-glass layer. The B2O3-concentration gradient across the borosilicate-glass, and the concomitant gradient in the ZrO2 solubility, appear to be responsible for this interesting phenomenon. Overall, this study has furthered our understanding of the inhomogeneous oxidation behavior of ZrB2SiC UHTC composites and its mitigation through new insights.  Acknowledgements  This research was supported by the U.S. Office of Naval Research (grant# N000141310459) monitored by Drs. L.T. Kabacoff and E.J. Wuchina. We thank Drs. T.A. Parthasarathy and A.R. Krause for fruitful discussions.  References  [4]  [11]  [13]  [14]  [8]  [9]  [10]  ceramics  in aerospace propulsion, Nat.  [1] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. Zhou, Ultra-high Temperature Ceramics: Materials for Extreme Environment Applications, John Wiley & Sons, Hoboken, NJ, 2014. [2] N.P. Padture, Advanced structural Mater. 15 (2016) 804e809. [3] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for extreme environments, Scr. Mater. 129 (2016) 94e99. T.A. Jackson, D.R. Eklund, A.J. Fink, High speed propulsion: performance advantage of advanced materials, J. Mater. Sci. 39 (2004) 5905e5913. 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Karlsdottir, J.W. Halloran, Formation of oxide scales on zirconium diborideesilicon carbide composites during oxidation: relation of subscale recession to liquid oxide flow, J. Am. Ceram. Soc. 91 (2008) 3652e3658. [18] K. Shugart, W. Jennings, E.J. Opila, Initial stages of ZrB2-30vol% SiC oxidation at 1500 \\x0e C, J. Am. Ceram. Soc. 97 (2014) 1645e1651. T.A. Parthasarathy, R.A. Rapp, M. Opeka, M.K. Cinibulk, Modeling oxidation kinetics of SiC-containing refractory diborides, J. Am. Ceram. Soc. 95 (2012) 338e349. [20] K. Shugart, S. Liu, F. Craven, E.J. Opila, Determination of retained B2O3 content in ZrB2-30 vol% SiC oxide scales, J. Am. Ceram. Soc. 98 (2015) 287e295. [21] R.H. Doremus, Handbook of Glass Properties, Academic Press, New York, 1986.  formation  [16]  [17]  [19]  \\x0c']"
},{
  "_id": 109,
  "PDF": "Initial oxidation behaviors of ZrB2-SiC-ZrC ternary composites above 2000 C.pdf",
  "Text": "[\"Journal of Alloys and Compounds 731 (2018) 310e317  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : h t t p : / / w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  Initial oxidation behaviors of ZrB2-SiC-ZrC ternary composites above 2000 \\x0eC  Ryo Inoue a, *, Yutaro Arai a, b, Yuki Kubota a, c, Yasuo Kogo a, Ken Goto a, d  a Department of Materials Science and Technology, Tokyo University of Science, 6-3-1 Nijyuku, Katsushika-ku, Tokyo, 125-8585, Japan b Department of Advanced Interdisciplinary Studies, Graduate School of Engineering, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo, 153-8904,  Japan c Japan Aerospace Exploration Agency (JAXA), Structures and Advanced Composite Research Unit, 6-13-1 Ohsawa, Mitaka-shi, Tokyo, 181-0015, d Department of Space Flight Systems,  Institute of Space and Astronautical Science (ISAS),  Japan Aerospace Exploration Agency (JAXA), 3-1-1 Yoshinodai,  Japan  Chuo-ku, Sagamihara, Kanagawa, 252-0222,  Japan  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 19 August 2017 Received in revised form 4 October 2017 Accepted 5 October 2017 Available online 6 October 2017  Keywords:  Ultra-high temperature ceramics (UHTCs) ZrB2 ZrC Resistance heating Oxidation  The initial oxidation behavior of ZrB2eSiCeZrC (ZSZ) composites above 2000 \\x0e C after the initial ~5e10 s was examined by an electric heating system. The morphology of the oxide layer was shown to depend on the ZrC content. Volume expansion during the conversion of ZrC to ZrO2 contributed to the morphology of the surface oxide layer. Above 2000 \\x0e C, the ZSZ composite with the highest ZrC amount had the lowest oxide layer thickness. This study clearly showed the efficacy of ZrC as an additive to significantly alter the oxide layer morphology relative to that observed for ZrB2eSiC composites tested under the same condition.  © 2017 Elsevier B.V. All rights reserved.  1.  Introduction  Ultra-high-temperature ceramics (UHTCs) are explored for use in nose cones and leading edges for hypersonic and re-entry vehicles. ZrB2 and its composites are a widely studied class of UHTCs. The oxidation of monolithic ZrB2 forms ZrO2 and B2O3. B2O3 acts as a component of the surface protective layer; however, it evaporates above 1200 \\x0e C [1e8]. SiC particles are considered effective additives to ZrB2 because the SiO2 formed by the oxidation of SiC protects unreacted regions. Simultaneously, excessive pores are formed under the surface in SiC particle-dispersed ZrB2 matrix (hereafter denoted ZS) composites in a wide temperature range via the preferential (active) oxidation of SiC because condensed SiO2 cannot form. Instead, gaseous SiO and CO are formed under the oxide layer mainly composed of SiO2 and ZrO2 by active oxidation because of the low O2 partial pressure relative to that at the surface [1,3e8]. The pore-rich porous layer is denoted the “SiC-depleted layer” [1,3e8].  * Corresponding author. inoue.ryo@rs.tus.ac.jp (R.  E-mail address:  Inoue).  https://doi.org/10.1016/j.jallcom.2017.10.034 0925-8388/© 2017 Elsevier B.V. All rights reserved.  The dispersion of SiC particles effectively enhances the oxidation resistance of ZrB2, because SiO2 layer formation on the surface prevents O2 diffusion toward the unoxidized region. The formation of the SiC-depleted layer under the oxide layer is a potential problem because of the material's lack of load-bearing capacity. This layer may lead to spallation and the delamination of the oxidized regions on the surface. Excessive pore formation in ZS composites should be prevented to improve their oxidation resistances. Attempts to add transition-metal compounds, such as WC, TaC, TaSi2, and TaB2, have been performed in order to solve the abovementioned problems. Opila et al. oxidized ZS composites with added TaSi2 and TaC at 1627 \\x0e C and 1927 \\x0eC, respectively, where the Zr in ZrO2 was substituted by Ta [9]. O2 diffusion through the oxide was decreased because of the decrease in the concentration of O vacancies in the ZrO2 skeleton by the addition of Ta. They reported an improvement in the oxidation resistance of the ZS composite by the addition of TaSi2. However, the liquid phase produced by Ta2O5 and Ta2O5$6ZrO2 caused a decrease in the oxidation resistance of the ZS composites at 1927 \\x0eC. Hu et al. added CrB2, HfB2, TaB, AlN, and La2O3 to ZS composites and oxidized them at 1800 \\x0e C for 1 h [10]. They concluded that, despite the use of additives intended to reduce oxidative weight gain, the weight gain was increased by the  \\x0c\", 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  311  addition of 10 vol% CrB2, 10 vol% TaB, 5 vol% AlN, and 5 vol% La2O3 compared to that of the unmodified ZS composites; no significant change occurred with the addition of 10 vol% HfB2. Recently, we reported that the addition of ZrC particles was effective in suppressing the formation of SiC-depleted layers in ZS composites at various temperature ranges [11,12]. Arai et al. reported that the SiC-depleted layer was formed by oxidation up to 1500 \\x0e C and that this probably caused the delamination of the oxide layer [11]. Kubota et al. conducted oxidation tests for ZrB2eSiCeZrC (ZSZ) composites with different ZrC contents at temperatures reaching 1700 in air and the low O2 partial pressure of 2.0 \\x02 103 Pa [12]. Although monolithic ZrC is oxidized drastically with the formation of ZrO2 because ZrO2 provides no protection, as it has a much higher oxygen permeability than that of SiO2 [13e16], they concluded that the partial pressure of O2 and the ZrC content were critical in controlling the formation of the SiC-depleted layer because the volume expansion during ZrC oxidation prevents pore formation. These experimental results suggested the potential use in hot sections up to 1700 \\x0eC. of ZSZ for structural applications However, the applicability of ZSZ ternary composites as structural components at high temperatures above 1700 remains  \\x0e C  \\x0eC  unknown. The oxidation behaviors have not yet been evaluated because the fabrication of large and complex-geometry specimens by sintering for arc-jet facilities is difficult. The objective of this work is to evaluate the effectiveness of ZrC particle additives in improving the oxidation resistance of ZS composites above 2000 \\x0eC. For evaluation, fast heating tests were performed using a resistance heating system for ZSZ composites with different ZrB2/ZrC ratios. The system could heat the specimen above 2000 \\x0e C at very high heating rates with no dynamic pressure. Exposure time at the maximum temperature was limited; however, data on the initial oxidation behaviors of ZSZ composites was obtained.  2.  Experimental procedure  2.1.  Sample preparation  ZSZ composites and ZS (ZrB2-20 vol% SiC, hereafter denoted ZS20) were fabricated by spark plasma sintering (SPS). The microstructures of the polished surfaces are shown in Fig. 1(a)-(e). Details of the fabrication procedure were reported in our previous works  Fig. 1. Typical microstructure of ZrB2-SiC-ZrC composites: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34, (d) ZSZ20, and (e) ZS20 [17].  \\x0c', '312  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  Table 1 Composition of ZSZ composites. For comparison purposes, compositions of ZrB2eSiC composites [17] are also indicated in the table.  Sample  VZrB2  (vol.%)  VSiC  (vol.%)  VZrC  (vol.%)  Porosity (%)  Density (g/cm3)  ZSZ  ZS  64 50 34 20 20  20 34 50 64 80  16 16 16 16 20  64 50 34 20  0.2 0.6 0.9 1.3 0.4  6.0 5.9 5.8 5.7 5.5  Fig. 2. Surface temperature history during the test.  [11,12,17]. Table 1 shows the compositions of the raw materials used in this study. A disk-shaped specimen with a diameter of ~15 mm and a thickness of ~3 mm was prepared. The powder compact was sintered by SPS at 1950 \\x0eC for 2 min under the uniaxial pressure of ~20 MPa. The relative density of all composites exceeded 95%. After sintering, the surface and cross-section of the specimen was polished using a diamond paste with a particle size of 1 mm. A rectangular parallelepiped specimen with dimensions of 5 \\x02 5 \\x02 3 mm (length \\x02 width \\x02 thickness) was cut from the as-sintered specimen using a diamond-impregnated blade.  2.2. Oxidation tests by resistance heating  Rapid oxidation tests above 2000 \\x0eC were performed using a specially designed electric heating system. Detailed conditions of the system have been described elsewhere [17]. The temperature  increase rate, maximum temperature, and holding time at the \\x0eC/s, \\x0e C, maximum temperature were set to ~300 ~2000 and ~5e10 s, respectively. From the maximum temperature, the specimen was cooled naturally. In the present study, we focused on the above 2000 \\x0eC initial oxidation behavior of ZSZ composites at without dynamic pressure and oxidation during heating. Thus, the holding time at the maximum temperature is lower than that in previous studies [11,12]. The surface temperature during the test was measured using a radiation thermometer. The emissivity was tentatively set at 1, because the value during oxidative reactions was initially unclear. Thus, the temperature measured by the radiation thermometer was lower than the actual specimen temperature. We ensured that the surface temperature of the specimen measured by the present system reached above 2000 \\x0eC because the maximum temperatures of all tests exceeded 2000 \\x0e C. After the test, the specimen was mounted in epoxy resin. The specimen was carefully cut and the cross-section was observed by scanning electron microscopy (SEM, JCM-6000plus, JEOL, Tokyo, Japan) with an accelerating voltage of 15 kV. Elemental analysis of the surface was also performed by energy dispersive X-ray spectroscopy (EDS, JED-2300, JEOL, Japan). Image analysis was also performed using captured images with a resolution of 1280 \\x02 1040 pixels using image-processing software (ImageJ, NIH, NA, USA). The accuracy of the measurements was ~0.03 mm.  3.  Experimental results  Fig. 2 shows the relationship between the surface temperature and test time during the oxidation tests. The surface temperature of the ZSZ composites increases at a rate of ~400 \\x0e C/s and reaches above 2000 \\x0e C. Fig. 3 shows general views of the oxidized ZSZ composites. ZSZ64 and ZSZ20 are shown as typical views of tested specimens. The images indicate that the microstructures evolved in oxidation testing appear to differ from those in un-oxidized regions; in addition, partial delamination occurs after the test. This demonstrates that the oxidation of ZSZ composites proceeds by the fast heating method used in the present study. Fig. 4 shows typical cross-sectional views after the oxidation of ZS20 [17]. ZS20 has a dense layer with the thickness of ~2e3 mm and a porous layer of ~10 mm in thickness with continuous pores with the diameter of ~1e2 mm. Typical cross-sectional views after the oxidation of ZSZ are also shown in Fig. 5. ZSZ20 has a partially dense layer near the surface and a porous layer containing continuous pores with the diameter of ~1e2 mm beneath the dense layer. ZSZ34 has a dense layer with the thickness of ~5 mm and a  Fig. 3. Typical SEM photographs of (a) ZSZ64 and (b) ZSZ20: General views of oxidized specimens.  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  313  detected. Meanwhile, Si and O are undetected in the ~2e5-mmthick region between the oxide layer and the non-oxide region for the ZS composites; this region is denoted the SiC-depleted layer. From these results, the oxide layers formed after oxidation for ZS and ZSZ composites are as follows: ZS20 has ZrO2 and SiC-depleted layers as oxide layers. ZSZ20 has a partially dense ZrO2 layer and a porous ZrO2 layer with continuous pores as the oxide layer. ZSZ34 and ZSZ50 have both dense and porous ZrO2 layers. ZSZ64 has a dense ZrO2 layer with some pores.  4. Discussion  4.1.  Surface temperature during test  In the present study, it is expected that the measured temperature during the oxidation of the ZS and ZSZ composites varies depending on the change of emissivity of the surface during oxidation because of the measurements by the radiation thermometer. As mentioned in Section 2.2, the emissivity is set to 1 for all oxidation tests. However, the emissivity of each specimen varies during heating because of the oxidation. From Fig. 6, it is implied that the emissivity of the ZSZ composites during oxidation changes within the range of emissivity of ZrO2 (~0.62e0.75) [18]. However, it is known that the emissivity of ZS composites varies within the range of ~0.8e0.95 [19,20]. We have already confirmed that the temperature changes from 2000 to ~2200 \\x0eC when the emissivity is varied from 1 to 0.62, as discussed elsewhere [17].  4.2. Oxidation behaviors of ZSZ and ZS composites up to 2000 \\x0e C  For the condition of oxidation in the present study, dynamic pressure is not applied to the specimens, whereas it is applied by torch testing or arc-jet testing. Moreover, the specimens reach 2000 \\x0eC with the heating rate of ~300 \\x0e C/s, which is ~3e4 orders of  Fig. 4. Typical SEM photographs of ZS20 composites after oxidation testing [17].  porous layer with the thickness of ~10e20 mm beneath the dense layer. Although ZSZ50 also has a dense layer with the thickness of ~10 mm and a porous layer beneath the dense layer, the number of pores in the porous layer is decreased in comparison with that in ZS20, ZSZ20, and ZSZ34. In addition, the dense and porous layers have different contrast (light gray and dark gray in Fig. 5 (b)). This difference of contrast is also shown in ZSZ34 and ZSZ20 (Fig. 5 (c) and (d)). Furthermore, ZSZ64 has a dense layer with the thickness of ~10e20 mm, featuring some pores with the diameter of ~5 mm; no distinct porous layer appears in ZSZ64. These results clearly show that increases in the amount of ZrC cause increased formation of dense layers during oxidation. Fig. 6 shows EDS analyses for Zr, Si, O, and C from the crosssection after the oxidation test. O is detected in all layers formed by oxidation in Figs. 4 and 5. The post-oxidation dense and porous layers are identified as oxides because O is detected in these layers; they are denoted “oxide layers.” In addition, the oxide layer comprises ZrO2 for both the ZSZ and ZS composites; only Zr and O are  Fig. 5. Typical SEM photographs of ZSZ composites after oxidation testing: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34 and (d) ZSZ20.  \\x0c', '314  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  Fig. 6. Elemental mapping of cross-sections for oxidized ZSZ and ZS composites.  magnitude greater than the rate in oxidation tests conducted by electric furnaces in air. Thus, it is considered that the oxidation behaviors of ZS and ZSZ composites at 2000 \\x0eC without the effect of dynamic pressure and oxidation during heating can be evaluated in comparison with the oxidation tests mentioned above because of the fast heating conducted in this study. Fig. 7 shows a schematic of the oxidation mechanism of the ZS composites oxidized by fast heating above 2000 \\x0e C [17]. After testing, a porous ZrO2 layer and SiC-depleted layer are observed in the oxidized region of the ZS composites (ZS20). This indicates that a SiC-depleted region was formed after oxidation at 2000 because gaseous SiO is formed by the preferential (active) oxidation of SiC instead of ZrB2 [17]. Meanwhile, Fig. 8 shows  schematic  \\x0e C  the  of  the  oxidation  Fig. 7. Schematic of oxidation mechanism in ZS20 composites [17].  mechanism of ZSZ composites above 2000 \\x0eC. In the oxidation of ZSZ composites, no SiC-depleted layer is formed after oxidation; the behavior differs completely from that of ZS composites [1,3e8] and from that of ZSZ oxidized below 2000 \\x0e C [11,12,21]. With the oxidation of transition metal carbides, gaseous products such as CO and CO2 are formed. As a result, the layer generally contains high amounts of pores [13]. However, for ZSZ64, the oxidized region is composed of a ZrO2 layer with few pores (hereafter denoted as the dense ZrO2 layer). For ZSZ50 and ZSZ34, it is clearly identified that both the dense ZrO2 layer and a ZrO2 layer with large amounts of pores (hereafter denoted as the porous ZrO2 layer) are formed. The morphologies of ZSZ20 are similar to those of ZSZ50 and ZSZ34. However, extensive pores are evolved and networked in ZSZ20. As a result, a porous layer is formed between the partially dense ZrO2 layer and the porous ZrO2 layer. The ZSZ composite with the highest content of ZrC (ZSZ64) has a dense ZrO2 layer. It is suggested that the sintering of ZrO2, mainly formed by the oxidation of ZrC, occurs during oxidation above 2000 \\x0eC. The volume expansion during oxidation of ZrC may, in addition to the high temperature, act as a driving force for sintering because the sintering of ZrO2-containing ceramics is promoted by volume expansion [22]. For ZS composites, ZrO2 is only formed by the oxidation of ZrB2; the volume expansion in the conversion from ZrB2 to ZrO2 is 16 vol%, lower than that from ZrC to ZrO2 [11,12,23e25]. Thus, it is implied that sintering of the oxide layer is limited for ZS composites during oxidation. However, the 33% volume expansion accompanying the conversion of ZrC to ZrO2 is twice that of ZrB2 to ZrO2. For smaller amounts of ZrC (for ZSZ50, 34, and 20), densification does not proceed because the volume expansion of ZrC is insufficient; therefore, the porous ZrO2 layer is observed. Although it is confirmed that increases in the amount of  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  315  Fig. 8. Schematic of oxidation mechanism in ZSZ composites: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34, and (d) ZSZ20.  ZrC cause the formation of the dense ZrO2 layer, further studies are required to prove the effect of volume expansion during oxidation by ZrC for the sintering of the oxide layer. Fig. 9 shows the relationship between the ZrC content and the average thickness of the oxide layer, hoxi. In the present study, hoxi is defined as hoxi ~ Shi/N, where hi is the thickness of the oxide layer measured by image analysis and N is the number of measurements made. From Fig. 9, the range of the thickness of ZS and ZSZ composites oxidized in the present study is ~20e30 mm and no effects of the difference of composition appear. This result is entirely different from the results for the oxidation of ZSZ composites below 2000 \\x0e C [11,12]. Although the thickness of the oxide layer in ZSZ64 is almost equal to that for ZS20 [17], a porous SiC-depleted layer is formed in ZS20; the thickness of the SiC-depleted layer is more than 50% of the thickness of the oxide layer formed during oxidation. Thus, the ZSZ composites have higher oxidation resistance because dense oxide layers are formed on the ZSZ composites in oxidation above 2000 \\x0e C, compared to the ZS composites.  4.3. Oxidation mechanism  As mentioned above, the oxidation and volume expansion of ZrC during oxidation are important factors in realizing oxidation resistance for ZSZ composites. In the present study, SiO2 is hardly detected in the ZrO2 formed during the oxidation of the ZSZ composites. This phenomenon, observed under the conditions established in the present study, is unique. Thus, a discussion in terms of thermodynamics is conducted as follows. The O2 partial pressure and vapor pressure of the reaction products dominate the oxidation reactions of ZrB2eSiC-based materials. Fig. 10 (a) and (b) show the volatility diagram for the oxidation of ZSZ composites below and above 2000 \\x0e C. Fig. 10 (c) shows an enlarged view of Fig. 10 (a) and (b). The thermodynamic  data used for calculation was obtained from the NIST-JANAF tables [26]. The volatility diagrams for the oxidation of ZS and ZSZ composites are discussed based on the following chemical reactions [11,12]:  ZrB2 (s) þ 5/2 O2 (g) / ZrO2 (s) þ B2O3 (l)  B2O3 (l) / B2O3 (g)  SiC (s) þ 3/2 O2 (g) / SiO2 (l) þ CO (g)  SiC (s) þ O2 (g) / SiO (g) þ CO (g)  (1)  (2)  (3)  (4)  Fig. 9. Plots of surface oxide layer thickness as a function of ZrC content.  \\x0c', '316  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  SiO2 (l) / SiO (g) þ 1/2 O2 (g)  ZrC (s) þ 3/2 O2 (g) / ZrO2 (s) þ CO (g)  (5)  (6)  \\x0014.15 Pa,  From Fig. 10, the O2 partial pressures for the oxidation of ZrB2, SiC, and ZrC at 1727 \\x0e C are 10 \\x008.09 Pa, 10 \\x009.16 Pa, and 10 respectively. However, at 2027 \\x0e C, \\x005.43 Pa, 10 \\x006.27 Pa, they are 10 \\x0010.5 Pa, respectively. This indicates that the oxidation of ZrC and 10 occurs preferentially before the oxidation of ZrB2 and SiC, both below and above 2000 \\x0e C. In the oxidation of ZSZ composites at 1700 \\x0eC, preferential oxidation of ZrC over ZrB2 and SiC was reported by Kubota et al. [12]. Moreover, Fig. 10(c) also indicates that the difference in O2 partial pressure required for the oxidation of compounds in ZSZ composites above 2000 \\x0e C is greater than that below 2000 \\x0eC. Thus, for the evolution of SiO2 by the oxidation of SiC in ZS and ZSZ composites, or for the re-oxidation of SiO formed by the active oxidation of SiC, more O2 is required for oxidation above 2000 \\x0eC than for that below 2000 \\x0e C. Then, as shown in Fig. 6, a porous ZrO2 layer is formed by the oxidation of ZrB2 and the SiC-depleted layer is formed by the active oxidation of SiC for ZS composites. For ZSZ composites, the amount of O2 required for the oxidation of ZrC is much lower than that needed for the active oxidation of SiC. Therefore, it is considered that the oxidation of ZrC occurs preferentially compared to that of SiC. In terms of the reaction kinetics, the reaction rate constant (k) for the oxidation of ZrB2, SiC, and ZrC is described as k ¼ A$exp(Ea/RT), where A, Ea, R, and T are the frequency factor, activation energy, gas constant, and temperature, respectively. The activation energy for the oxidation of ZrB2 with the formation of B2O3 scale below 1100 \\x0e C is 83 kJ/mol; that with the evaporation of B2O3 above 1100 \\x0eC is 323 kJ/mol [12,27]. In addition, the activation energies for the oxidation of SiC (Reaction (4)) and ZrC (Reaction (6)) are 97 kJ/mol and 70 kJ/mol, respectively [28e30]. Thus, the oxidation of ZrC occurs preferentially compared to that of ZrB2 and SiC in terms of the reaction kinetics. Moreover, Reaction (5) is required for the evolution of SiO2 based on the volatility diagram (Fig. 10). Therefore, it is suggested that ZrC is oxidized preferentially over SiC, because the frequency factor for the reaction of solids tends to be higher than that for gas. From the above, the dense ZrO2 layer is formed by the oxidation of ZrC; SiO2 is not formed under the conditions in the present study. However, in the present study, the cooling time is longer than the holding time at the maximum temperature. From Fig. 10 (a) and (b), the minimum oxygen partial pressure to form SiO2 by reaction (3) \\x006.27 (2027 \\x0eC) to decreases with decreasing temperature (from 10 \\x009.16 (1727 \\x0e C)). Thus, SiO2 can be formed during cooling and the 10 difference of contrast in oxide layer mentioned in Section 3 (Fig. 5) is probably caused by SiO2 formed during oxidation. It is clear that a dense oxide layer is formed by the addition of ZrC during initial oxidation, while porous SiC-depleted layers do not form in ZSZ composites oxidized under same condition. Generally, transition metal carbides (ZrC) and SiC form gaseous species such as CO and SiO during oxidation and these gaseous species lead to formation of porous oxidized layer [13,14]. To further understand the effectiveness of the ZrC additive, the developed heating system should be extended to a much longer test duration to examine the effect of these gas species on ZrO2 layer. However, for the first time, to our best knowledge, we have identified that the effect of the addition of ZrC to ZS composite to promote the formation of a dense oxide layer appears only for oxidation above 2000 \\x0eC.  5.  Conclusions  In the present study,  rapid oxidation tests of ZSZ composites  Fig. 10. The volatility diagram for oxidation of ZSZ composites: T ¼ 2027 \\x0e C, and (c) enlarged view of (a) and (b).  (a) T ¼ 1727 \\x0e C,  (b)  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  317  were performed using an electric heating system. Microstructural observations and thermodynamics-based analysis were performed and the following conclusions were reached:  1.  In the present test conditions, no SiC-depleted layer formed after the oxidation of the ZSZ composites; instead, a dense ZrO2 layer formed on the ZSZ composite with the highest ZrC content. This indicated that the addition of ZrC was effective for improving the oxidation resistance of ZS composites. 2. Dense ZrO2 layers formed on ZSZ composites only after oxidation above 2000 \\x0e C. Meanwhile, the porosity of the ZrO2 layer increased with decreasing ZrC content in the ZSZ composites. 3. From the volatility diagrams for the oxidation of ZSZ composites, ZrC was concluded to oxidize preferentially over ZrB2 and SiC, even above 2000 \\x0e C. The disappearance of the SiC-depleted layer was caused by the oxidation of ZrC in the SiC-depleted layer, which then became a porous ZrO2 layer. 4. Thermodynamic analysis suggested that the formation of SiO2 is limited during the oxidation of ZSZ composites above 2000 \\x0e C because of the increased O2 partial pressure required to form SiO2 and because the activation energy for the oxidation of ZrC was higher than that for ZrB2 and SiC, unlike the oxidation of ZSZ composites below 2000 \\x0e C. However, because of the longer cooling time relative to the holding time at the maximum temperature, SiO2 is formed during cooling and it remains in oxide layer. 5. The effect of the addition of ZrC to ZrB2eSiC could be evaluated only by oxidation tests above 2000 \\x0e C. The fast heating method proposed in the present study was effective for evaluating the oxidation mechanism of ZSZ composites above 2000 \\x0e C.  References  [4]  [3]  formation  [1] W.G. Fahrenholtz, Thermodynamic analysis of ZrB2-SiC oxidation: of a SiC-depleted region, J. Am. Ceram. Soc. 90 (2007) 143e148. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347e1364. F. Monteverde, R. Savino, M. De Stefano Fumo, Dynamic oxidation of ultrahigh temperature ZrB2-SiC under high enthalpy supersonic flows, Corros. Sci. 53 (2011) 922e929. F. Monteverde, R. Savino, M. De Stefano Fumo, A. Di Maso, Plasma wind tunnel testing of ultra-high temperature ZrB2-SiC composites under hypersonic reentry conditions, J. Eur. Ceram. Soc. 30 (2010) 2313e2321. J. Han, P. Hu, X. Zhang, S. Meng, W. Han, Oxidation-resistant ZrB2-SiC composites at 2200\\x0e C, Compos. Sci. Technol. 68 (2008) 799e806. S.N. Karlsdottir, J.W. Halloran, Rapid oxidation characterization of ultra-high temperature ceramics, J. Am. Ceram. Soc. 90 (2007) 3233e3238. [7] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the  [5]  [6]  [8]  [9]  [15]  [10]  J. Eur.  and HfC,  oxidation of zirconium diborideesilicon carbide in air up to 1500\\x0e C, Ceram. Soc. 27 (2007) 2495e2501. F. Monteverde, R. Savino, Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions, J. Eur. Ceram. Soc. 27 (2007) 4797e4805. E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2and HfB2-based ultra-high temperature ceramics: effect of Ta additions, J. Mater. Sci. 39 (2004) 5969e5977. P. Hu, X.-H. Zhang, J.-C. Han, X.-G. Luo, S.-Y. Du, Effect of various additives on the oxidation behavior of ZrB2-based ultra-high-temperature ceramics at 1800\\x0e C, J. Am. Ceram. Soc. 93 (2010) 345e349. [11] Y. Arai, R. Inoue, H. Tanaka, Y. Kogo, K. Goto, In-situ observation of oxidation behavior in ZrB2-SiC-ZrC ternary composites up to 1500\\x0e C using hightemperature observation system, J. Ceram. Soc. Jpn. 124 (2016) 890e897. [12] Y. Kubota, H. Tanaka, Y. Arai, R. Inoue, Y. Kogo, K. Goto, Oxidation behavior of ZrB2-SiC-ZrC at 1700\\x0e C, J. Eur. Ceram. Soc. 37 (2017) 1187e1194. [13] R.F. Vojtovich, E.A. Pugach, High-temperature oxidation of ZrC Powder Metall. Met. Ceram. 12 (1973) 916e921. [14] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000\\x0e C hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887e5904. J. Doychak, T. Grobstein, The oxidation of high-temperature intermetallics, JOM 41 (1989) 30e31. [16] A.W. Smith, F.W. Meszaros, C.D. Amata, Permeability of zirconia, hafnia, and thoria to oxygen, J. Am. Ceram. Soc. 49 (1966) 240e244. [17] R. Inoue, Y. Arai, Y. Kubota, Oxidation behaviors of ZrB2-SiC binary composites above 2000\\x0e C, Ceram. Int. 43 (2017) 8081e8088. [18] D. Cubicciotti, The melting pointdcomposition diagram of niumdoxygen system, J. Am. Chem. Soc. 73 (1951) 2032e2035. J. Marschall, D.A. Pejakovi\\x13c, W.G. Fahrenholtz, G.E. Hilmas, S. Zhu, J. Ridge, D.G. Fletcher, C.O. Asma, J. Th€ome, Oxidation of ZrB2-SiC ultrahigh-temperature ceramic composites in dissociated air, J. Thermophys. Heat Transfer 23 (2009) 267e278. L. Scatteia, D. Alfano, F. Monteverde, J.-L. Sans, M. Balat-Pichelin, Effect of the machining method on the catalycity and emissivity of ZrB2 and ZrB2-HfB2based ceramics, J. Am. Ceram. Soc. 91 (2008) 1461e1468. I. Akin, G. Goller, Mechanical and oxidation behavior of spark plasma sintered ZrB2-ZrC-SiC composites, J. Ceram. Soc. Jpn. 120 (2012) 143e149. [22] M. Chen, C. Lu, J. Yu, Improvement in performance of MgOeCaO refractories by addition of nano-sized ZrO2, J. Eur. Ceram. Soc. 27 (2007) 4633e4638. Z. Wang, Z. Wu, G. Shi, The oxidation behaviors of a ZrB2-SiC-ZrC ceramic, Solid State Sci. 13 (2011) 534e538. Z. Wang, P. Zhou, Z. Wu, Effect of surface oxidation on thermal shock resistance of ZrB2eSiCeZrC ceramic at temperature difference from 800 to 1900\\x0e C, Corros. Sci. 98 (2015) 233e239. [25] H.-L. Liu, J.-X. Liu, H.-T. Liu, G.-J. Zhang, Changed oxidation behavior of ZrB2SiC ceramics with the addition of ZrC, Ceram. Int. 41 (2015) 8247e8251. [26] M.W. Chase Jr., NIST-JANAF Thermochemical Tables, fourth ed., American Institute of Physics, Woodbury, NY, 1998. J.B. Berkowitz-Mattuck, High-temperature oxidation III. Zirconium and hafnium diborides, J. Electrochem. Soc. 113 (9) (1966) 908e914. [28] A.K. Kuriakose, J.L. Margrave, The oxidation kinetics of zirconium diboride and zirconium carbide at high temperatures, J. Elec. Soc. 111 (1964) 827e831. J.W. Hinze, H.C. Graham, The active oxidation of Si and SiC in the viscous gasflow regime, J. Electrochem. Soc. 123 (1976) 1066e1073. [30] Y. Kubota, H. Hatta, T. Yoshinaka, T. Goto, T. Rong, Use of volume element methods to understand experimental differences in active/passive transitions and active oxidation rates for SiC, J. Am. Ceram. Soc. 96 (2013) 1317e1323.  zirco [23]  [24]  [20]  [21]  the  [19]  [27]  [29]  \\x0c']"
},{
  "_id": 110,
  "PDF": "Initial oxidation behaviors of ZrB2-SiC-ZrC ternary composites above 2000 °C.pdf",
  "Text": "[\"Journal of Alloys and Compounds 731 (2018) 310e317  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : h t t p : / / w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  Initial oxidation behaviors of ZrB2-SiC-ZrC ternary composites above 2000 \\x0eC  Ryo Inoue a, *, Yutaro Arai a, b, Yuki Kubota a, c, Yasuo Kogo a, Ken Goto a, d  a Department of Materials Science and Technology, Tokyo University of Science, 6-3-1 Nijyuku, Katsushika-ku, Tokyo, 125-8585, Japan b Department of Advanced Interdisciplinary Studies, Graduate School of Engineering, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo, 153-8904,  Japan c Japan Aerospace Exploration Agency (JAXA), Structures and Advanced Composite Research Unit, 6-13-1 Ohsawa, Mitaka-shi, Tokyo, 181-0015, d Department of Space Flight Systems,  Institute of Space and Astronautical Science (ISAS),  Japan Aerospace Exploration Agency (JAXA), 3-1-1 Yoshinodai,  Japan  Chuo-ku, Sagamihara, Kanagawa, 252-0222,  Japan  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 19 August 2017 Received in revised form 4 October 2017 Accepted 5 October 2017 Available online 6 October 2017  Keywords:  Ultra-high temperature ceramics (UHTCs) ZrB2 ZrC Resistance heating Oxidation  The initial oxidation behavior of ZrB2eSiCeZrC (ZSZ) composites above 2000 \\x0e C after the initial ~5e10 s was examined by an electric heating system. The morphology of the oxide layer was shown to depend on the ZrC content. Volume expansion during the conversion of ZrC to ZrO2 contributed to the morphology of the surface oxide layer. Above 2000 \\x0e C, the ZSZ composite with the highest ZrC amount had the lowest oxide layer thickness. This study clearly showed the efficacy of ZrC as an additive to significantly alter the oxide layer morphology relative to that observed for ZrB2eSiC composites tested under the same condition.  © 2017 Elsevier B.V. All rights reserved.  1.  Introduction  Ultra-high-temperature ceramics (UHTCs) are explored for use in nose cones and leading edges for hypersonic and re-entry vehicles. ZrB2 and its composites are a widely studied class of UHTCs. The oxidation of monolithic ZrB2 forms ZrO2 and B2O3. B2O3 acts as a component of the surface protective layer; however, it evaporates above 1200 \\x0e C [1e8]. SiC particles are considered effective additives to ZrB2 because the SiO2 formed by the oxidation of SiC protects unreacted regions. Simultaneously, excessive pores are formed under the surface in SiC particle-dispersed ZrB2 matrix (hereafter denoted ZS) composites in a wide temperature range via the preferential (active) oxidation of SiC because condensed SiO2 cannot form. Instead, gaseous SiO and CO are formed under the oxide layer mainly composed of SiO2 and ZrO2 by active oxidation because of the low O2 partial pressure relative to that at the surface [1,3e8]. The pore-rich porous layer is denoted the “SiC-depleted layer” [1,3e8].  * Corresponding author. inoue.ryo@rs.tus.ac.jp (R.  E-mail address:  Inoue).  https://doi.org/10.1016/j.jallcom.2017.10.034 0925-8388/© 2017 Elsevier B.V. All rights reserved.  The dispersion of SiC particles effectively enhances the oxidation resistance of ZrB2, because SiO2 layer formation on the surface prevents O2 diffusion toward the unoxidized region. The formation of the SiC-depleted layer under the oxide layer is a potential problem because of the material's lack of load-bearing capacity. This layer may lead to spallation and the delamination of the oxidized regions on the surface. Excessive pore formation in ZS composites should be prevented to improve their oxidation resistances. Attempts to add transition-metal compounds, such as WC, TaC, TaSi2, and TaB2, have been performed in order to solve the abovementioned problems. Opila et al. oxidized ZS composites with added TaSi2 and TaC at 1627 \\x0e C and 1927 \\x0eC, respectively, where the Zr in ZrO2 was substituted by Ta [9]. O2 diffusion through the oxide was decreased because of the decrease in the concentration of O vacancies in the ZrO2 skeleton by the addition of Ta. They reported an improvement in the oxidation resistance of the ZS composite by the addition of TaSi2. However, the liquid phase produced by Ta2O5 and Ta2O5$6ZrO2 caused a decrease in the oxidation resistance of the ZS composites at 1927 \\x0eC. Hu et al. added CrB2, HfB2, TaB, AlN, and La2O3 to ZS composites and oxidized them at 1800 \\x0e C for 1 h [10]. They concluded that, despite the use of additives intended to reduce oxidative weight gain, the weight gain was increased by the  \\x0c\", 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  311  addition of 10 vol% CrB2, 10 vol% TaB, 5 vol% AlN, and 5 vol% La2O3 compared to that of the unmodified ZS composites; no significant change occurred with the addition of 10 vol% HfB2. Recently, we reported that the addition of ZrC particles was effective in suppressing the formation of SiC-depleted layers in ZS composites at various temperature ranges [11,12]. Arai et al. reported that the SiC-depleted layer was formed by oxidation up to 1500 \\x0e C and that this probably caused the delamination of the oxide layer [11]. Kubota et al. conducted oxidation tests for ZrB2eSiCeZrC (ZSZ) composites with different ZrC contents at temperatures reaching 1700 in air and the low O2 partial pressure of 2.0 \\x02 103 Pa [12]. Although monolithic ZrC is oxidized drastically with the formation of ZrO2 because ZrO2 provides no protection, as it has a much higher oxygen permeability than that of SiO2 [13e16], they concluded that the partial pressure of O2 and the ZrC content were critical in controlling the formation of the SiC-depleted layer because the volume expansion during ZrC oxidation prevents pore formation. These experimental results suggested the potential use in hot sections up to 1700 \\x0eC. of ZSZ for structural applications However, the applicability of ZSZ ternary composites as structural components at high temperatures above 1700 remains  \\x0e C  \\x0eC  unknown. The oxidation behaviors have not yet been evaluated because the fabrication of large and complex-geometry specimens by sintering for arc-jet facilities is difficult. The objective of this work is to evaluate the effectiveness of ZrC particle additives in improving the oxidation resistance of ZS composites above 2000 \\x0eC. For evaluation, fast heating tests were performed using a resistance heating system for ZSZ composites with different ZrB2/ZrC ratios. The system could heat the specimen above 2000 \\x0e C at very high heating rates with no dynamic pressure. Exposure time at the maximum temperature was limited; however, data on the initial oxidation behaviors of ZSZ composites was obtained.  2.  Experimental procedure  2.1.  Sample preparation  ZSZ composites and ZS (ZrB2-20 vol% SiC, hereafter denoted ZS20) were fabricated by spark plasma sintering (SPS). The microstructures of the polished surfaces are shown in Fig. 1(a)-(e). Details of the fabrication procedure were reported in our previous works  Fig. 1. Typical microstructure of ZrB2-SiC-ZrC composites: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34, (d) ZSZ20, and (e) ZS20 [17].  \\x0c', '312  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  Table 1 Composition of ZSZ composites. For comparison purposes, compositions of ZrB2eSiC composites [17] are also indicated in the table.  Sample  VZrB2  (vol.%)  VSiC  (vol.%)  VZrC  (vol.%)  Porosity (%)  Density (g/cm3)  ZSZ  ZS  64 50 34 20 20  20 34 50 64 80  16 16 16 16 20  64 50 34 20  0.2 0.6 0.9 1.3 0.4  6.0 5.9 5.8 5.7 5.5  Fig. 2. Surface temperature history during the test.  [11,12,17]. Table 1 shows the compositions of the raw materials used in this study. A disk-shaped specimen with a diameter of ~15 mm and a thickness of ~3 mm was prepared. The powder compact was sintered by SPS at 1950 \\x0eC for 2 min under the uniaxial pressure of ~20 MPa. The relative density of all composites exceeded 95%. After sintering, the surface and cross-section of the specimen was polished using a diamond paste with a particle size of 1 mm. A rectangular parallelepiped specimen with dimensions of 5 \\x02 5 \\x02 3 mm (length \\x02 width \\x02 thickness) was cut from the as-sintered specimen using a diamond-impregnated blade.  2.2. Oxidation tests by resistance heating  Rapid oxidation tests above 2000 \\x0eC were performed using a specially designed electric heating system. Detailed conditions of the system have been described elsewhere [17]. The temperature  increase rate, maximum temperature, and holding time at the \\x0eC/s, \\x0e C, maximum temperature were set to ~300 ~2000 and ~5e10 s, respectively. From the maximum temperature, the specimen was cooled naturally. In the present study, we focused on the above 2000 \\x0eC initial oxidation behavior of ZSZ composites at without dynamic pressure and oxidation during heating. Thus, the holding time at the maximum temperature is lower than that in previous studies [11,12]. The surface temperature during the test was measured using a radiation thermometer. The emissivity was tentatively set at 1, because the value during oxidative reactions was initially unclear. Thus, the temperature measured by the radiation thermometer was lower than the actual specimen temperature. We ensured that the surface temperature of the specimen measured by the present system reached above 2000 \\x0eC because the maximum temperatures of all tests exceeded 2000 \\x0e C. After the test, the specimen was mounted in epoxy resin. The specimen was carefully cut and the cross-section was observed by scanning electron microscopy (SEM, JCM-6000plus, JEOL, Tokyo, Japan) with an accelerating voltage of 15 kV. Elemental analysis of the surface was also performed by energy dispersive X-ray spectroscopy (EDS, JED-2300, JEOL, Japan). Image analysis was also performed using captured images with a resolution of 1280 \\x02 1040 pixels using image-processing software (ImageJ, NIH, NA, USA). The accuracy of the measurements was ~0.03 mm.  3.  Experimental results  Fig. 2 shows the relationship between the surface temperature and test time during the oxidation tests. The surface temperature of the ZSZ composites increases at a rate of ~400 \\x0e C/s and reaches above 2000 \\x0e C. Fig. 3 shows general views of the oxidized ZSZ composites. ZSZ64 and ZSZ20 are shown as typical views of tested specimens. The images indicate that the microstructures evolved in oxidation testing appear to differ from those in un-oxidized regions; in addition, partial delamination occurs after the test. This demonstrates that the oxidation of ZSZ composites proceeds by the fast heating method used in the present study. Fig. 4 shows typical cross-sectional views after the oxidation of ZS20 [17]. ZS20 has a dense layer with the thickness of ~2e3 mm and a porous layer of ~10 mm in thickness with continuous pores with the diameter of ~1e2 mm. Typical cross-sectional views after the oxidation of ZSZ are also shown in Fig. 5. ZSZ20 has a partially dense layer near the surface and a porous layer containing continuous pores with the diameter of ~1e2 mm beneath the dense layer. ZSZ34 has a dense layer with the thickness of ~5 mm and a  Fig. 3. Typical SEM photographs of (a) ZSZ64 and (b) ZSZ20: General views of oxidized specimens.  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  313  detected. Meanwhile, Si and O are undetected in the ~2e5-mmthick region between the oxide layer and the non-oxide region for the ZS composites; this region is denoted the SiC-depleted layer. From these results, the oxide layers formed after oxidation for ZS and ZSZ composites are as follows: ZS20 has ZrO2 and SiC-depleted layers as oxide layers. ZSZ20 has a partially dense ZrO2 layer and a porous ZrO2 layer with continuous pores as the oxide layer. ZSZ34 and ZSZ50 have both dense and porous ZrO2 layers. ZSZ64 has a dense ZrO2 layer with some pores.  4. Discussion  4.1.  Surface temperature during test  In the present study, it is expected that the measured temperature during the oxidation of the ZS and ZSZ composites varies depending on the change of emissivity of the surface during oxidation because of the measurements by the radiation thermometer. As mentioned in Section 2.2, the emissivity is set to 1 for all oxidation tests. However, the emissivity of each specimen varies during heating because of the oxidation. From Fig. 6, it is implied that the emissivity of the ZSZ composites during oxidation changes within the range of emissivity of ZrO2 (~0.62e0.75) [18]. However, it is known that the emissivity of ZS composites varies within the range of ~0.8e0.95 [19,20]. We have already confirmed that the temperature changes from 2000 to ~2200 \\x0eC when the emissivity is varied from 1 to 0.62, as discussed elsewhere [17].  4.2. Oxidation behaviors of ZSZ and ZS composites up to 2000 \\x0e C  For the condition of oxidation in the present study, dynamic pressure is not applied to the specimens, whereas it is applied by torch testing or arc-jet testing. Moreover, the specimens reach 2000 \\x0eC with the heating rate of ~300 \\x0e C/s, which is ~3e4 orders of  Fig. 4. Typical SEM photographs of ZS20 composites after oxidation testing [17].  porous layer with the thickness of ~10e20 mm beneath the dense layer. Although ZSZ50 also has a dense layer with the thickness of ~10 mm and a porous layer beneath the dense layer, the number of pores in the porous layer is decreased in comparison with that in ZS20, ZSZ20, and ZSZ34. In addition, the dense and porous layers have different contrast (light gray and dark gray in Fig. 5 (b)). This difference of contrast is also shown in ZSZ34 and ZSZ20 (Fig. 5 (c) and (d)). Furthermore, ZSZ64 has a dense layer with the thickness of ~10e20 mm, featuring some pores with the diameter of ~5 mm; no distinct porous layer appears in ZSZ64. These results clearly show that increases in the amount of ZrC cause increased formation of dense layers during oxidation. Fig. 6 shows EDS analyses for Zr, Si, O, and C from the crosssection after the oxidation test. O is detected in all layers formed by oxidation in Figs. 4 and 5. The post-oxidation dense and porous layers are identified as oxides because O is detected in these layers; they are denoted “oxide layers.” In addition, the oxide layer comprises ZrO2 for both the ZSZ and ZS composites; only Zr and O are  Fig. 5. Typical SEM photographs of ZSZ composites after oxidation testing: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34 and (d) ZSZ20.  \\x0c', '314  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  Fig. 6. Elemental mapping of cross-sections for oxidized ZSZ and ZS composites.  magnitude greater than the rate in oxidation tests conducted by electric furnaces in air. Thus, it is considered that the oxidation behaviors of ZS and ZSZ composites at 2000 \\x0eC without the effect of dynamic pressure and oxidation during heating can be evaluated in comparison with the oxidation tests mentioned above because of the fast heating conducted in this study. Fig. 7 shows a schematic of the oxidation mechanism of the ZS composites oxidized by fast heating above 2000 \\x0e C [17]. After testing, a porous ZrO2 layer and SiC-depleted layer are observed in the oxidized region of the ZS composites (ZS20). This indicates that a SiC-depleted region was formed after oxidation at 2000 because gaseous SiO is formed by the preferential (active) oxidation of SiC instead of ZrB2 [17]. Meanwhile, Fig. 8 shows  schematic  \\x0e C  the  of  the  oxidation  Fig. 7. Schematic of oxidation mechanism in ZS20 composites [17].  mechanism of ZSZ composites above 2000 \\x0eC. In the oxidation of ZSZ composites, no SiC-depleted layer is formed after oxidation; the behavior differs completely from that of ZS composites [1,3e8] and from that of ZSZ oxidized below 2000 \\x0e C [11,12,21]. With the oxidation of transition metal carbides, gaseous products such as CO and CO2 are formed. As a result, the layer generally contains high amounts of pores [13]. However, for ZSZ64, the oxidized region is composed of a ZrO2 layer with few pores (hereafter denoted as the dense ZrO2 layer). For ZSZ50 and ZSZ34, it is clearly identified that both the dense ZrO2 layer and a ZrO2 layer with large amounts of pores (hereafter denoted as the porous ZrO2 layer) are formed. The morphologies of ZSZ20 are similar to those of ZSZ50 and ZSZ34. However, extensive pores are evolved and networked in ZSZ20. As a result, a porous layer is formed between the partially dense ZrO2 layer and the porous ZrO2 layer. The ZSZ composite with the highest content of ZrC (ZSZ64) has a dense ZrO2 layer. It is suggested that the sintering of ZrO2, mainly formed by the oxidation of ZrC, occurs during oxidation above 2000 \\x0eC. The volume expansion during oxidation of ZrC may, in addition to the high temperature, act as a driving force for sintering because the sintering of ZrO2-containing ceramics is promoted by volume expansion [22]. For ZS composites, ZrO2 is only formed by the oxidation of ZrB2; the volume expansion in the conversion from ZrB2 to ZrO2 is 16 vol%, lower than that from ZrC to ZrO2 [11,12,23e25]. Thus, it is implied that sintering of the oxide layer is limited for ZS composites during oxidation. However, the 33% volume expansion accompanying the conversion of ZrC to ZrO2 is twice that of ZrB2 to ZrO2. For smaller amounts of ZrC (for ZSZ50, 34, and 20), densification does not proceed because the volume expansion of ZrC is insufficient; therefore, the porous ZrO2 layer is observed. Although it is confirmed that increases in the amount of  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  315  Fig. 8. Schematic of oxidation mechanism in ZSZ composites: (a) ZSZ64, (b) ZSZ50, (c) ZSZ34, and (d) ZSZ20.  ZrC cause the formation of the dense ZrO2 layer, further studies are required to prove the effect of volume expansion during oxidation by ZrC for the sintering of the oxide layer. Fig. 9 shows the relationship between the ZrC content and the average thickness of the oxide layer, hoxi. In the present study, hoxi is defined as hoxi ~ Shi/N, where hi is the thickness of the oxide layer measured by image analysis and N is the number of measurements made. From Fig. 9, the range of the thickness of ZS and ZSZ composites oxidized in the present study is ~20e30 mm and no effects of the difference of composition appear. This result is entirely different from the results for the oxidation of ZSZ composites below 2000 \\x0e C [11,12]. Although the thickness of the oxide layer in ZSZ64 is almost equal to that for ZS20 [17], a porous SiC-depleted layer is formed in ZS20; the thickness of the SiC-depleted layer is more than 50% of the thickness of the oxide layer formed during oxidation. Thus, the ZSZ composites have higher oxidation resistance because dense oxide layers are formed on the ZSZ composites in oxidation above 2000 \\x0e C, compared to the ZS composites.  4.3. Oxidation mechanism  As mentioned above, the oxidation and volume expansion of ZrC during oxidation are important factors in realizing oxidation resistance for ZSZ composites. In the present study, SiO2 is hardly detected in the ZrO2 formed during the oxidation of the ZSZ composites. This phenomenon, observed under the conditions established in the present study, is unique. Thus, a discussion in terms of thermodynamics is conducted as follows. The O2 partial pressure and vapor pressure of the reaction products dominate the oxidation reactions of ZrB2eSiC-based materials. Fig. 10 (a) and (b) show the volatility diagram for the oxidation of ZSZ composites below and above 2000 \\x0e C. Fig. 10 (c) shows an enlarged view of Fig. 10 (a) and (b). The thermodynamic  data used for calculation was obtained from the NIST-JANAF tables [26]. The volatility diagrams for the oxidation of ZS and ZSZ composites are discussed based on the following chemical reactions [11,12]:  ZrB2 (s) þ 5/2 O2 (g) / ZrO2 (s) þ B2O3 (l)  B2O3 (l) / B2O3 (g)  SiC (s) þ 3/2 O2 (g) / SiO2 (l) þ CO (g)  SiC (s) þ O2 (g) / SiO (g) þ CO (g)  (1)  (2)  (3)  (4)  Fig. 9. Plots of surface oxide layer thickness as a function of ZrC content.  \\x0c', '316  R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  SiO2 (l) / SiO (g) þ 1/2 O2 (g)  ZrC (s) þ 3/2 O2 (g) / ZrO2 (s) þ CO (g)  (5)  (6)  \\x0014.15 Pa,  From Fig. 10, the O2 partial pressures for the oxidation of ZrB2, SiC, and ZrC at 1727 \\x0e C are 10 \\x008.09 Pa, 10 \\x009.16 Pa, and 10 respectively. However, at 2027 \\x0e C, \\x005.43 Pa, 10 \\x006.27 Pa, they are 10 \\x0010.5 Pa, respectively. This indicates that the oxidation of ZrC and 10 occurs preferentially before the oxidation of ZrB2 and SiC, both below and above 2000 \\x0e C. In the oxidation of ZSZ composites at 1700 \\x0eC, preferential oxidation of ZrC over ZrB2 and SiC was reported by Kubota et al. [12]. Moreover, Fig. 10(c) also indicates that the difference in O2 partial pressure required for the oxidation of compounds in ZSZ composites above 2000 \\x0e C is greater than that below 2000 \\x0eC. Thus, for the evolution of SiO2 by the oxidation of SiC in ZS and ZSZ composites, or for the re-oxidation of SiO formed by the active oxidation of SiC, more O2 is required for oxidation above 2000 \\x0eC than for that below 2000 \\x0e C. Then, as shown in Fig. 6, a porous ZrO2 layer is formed by the oxidation of ZrB2 and the SiC-depleted layer is formed by the active oxidation of SiC for ZS composites. For ZSZ composites, the amount of O2 required for the oxidation of ZrC is much lower than that needed for the active oxidation of SiC. Therefore, it is considered that the oxidation of ZrC occurs preferentially compared to that of SiC. In terms of the reaction kinetics, the reaction rate constant (k) for the oxidation of ZrB2, SiC, and ZrC is described as k ¼ A$exp(Ea/RT), where A, Ea, R, and T are the frequency factor, activation energy, gas constant, and temperature, respectively. The activation energy for the oxidation of ZrB2 with the formation of B2O3 scale below 1100 \\x0e C is 83 kJ/mol; that with the evaporation of B2O3 above 1100 \\x0eC is 323 kJ/mol [12,27]. In addition, the activation energies for the oxidation of SiC (Reaction (4)) and ZrC (Reaction (6)) are 97 kJ/mol and 70 kJ/mol, respectively [28e30]. Thus, the oxidation of ZrC occurs preferentially compared to that of ZrB2 and SiC in terms of the reaction kinetics. Moreover, Reaction (5) is required for the evolution of SiO2 based on the volatility diagram (Fig. 10). Therefore, it is suggested that ZrC is oxidized preferentially over SiC, because the frequency factor for the reaction of solids tends to be higher than that for gas. From the above, the dense ZrO2 layer is formed by the oxidation of ZrC; SiO2 is not formed under the conditions in the present study. However, in the present study, the cooling time is longer than the holding time at the maximum temperature. From Fig. 10 (a) and (b), the minimum oxygen partial pressure to form SiO2 by reaction (3) \\x006.27 (2027 \\x0eC) to decreases with decreasing temperature (from 10 \\x009.16 (1727 \\x0e C)). Thus, SiO2 can be formed during cooling and the 10 difference of contrast in oxide layer mentioned in Section 3 (Fig. 5) is probably caused by SiO2 formed during oxidation. It is clear that a dense oxide layer is formed by the addition of ZrC during initial oxidation, while porous SiC-depleted layers do not form in ZSZ composites oxidized under same condition. Generally, transition metal carbides (ZrC) and SiC form gaseous species such as CO and SiO during oxidation and these gaseous species lead to formation of porous oxidized layer [13,14]. To further understand the effectiveness of the ZrC additive, the developed heating system should be extended to a much longer test duration to examine the effect of these gas species on ZrO2 layer. However, for the first time, to our best knowledge, we have identified that the effect of the addition of ZrC to ZS composite to promote the formation of a dense oxide layer appears only for oxidation above 2000 \\x0eC.  5.  Conclusions  In the present study,  rapid oxidation tests of ZSZ composites  Fig. 10. The volatility diagram for oxidation of ZSZ composites: T ¼ 2027 \\x0e C, and (c) enlarged view of (a) and (b).  (a) T ¼ 1727 \\x0e C,  (b)  \\x0c', 'R.  Inoue et al.  /  Journal of Alloys and Compounds 731 (2018) 310e317  317  were performed using an electric heating system. Microstructural observations and thermodynamics-based analysis were performed and the following conclusions were reached:  1.  In the present test conditions, no SiC-depleted layer formed after the oxidation of the ZSZ composites; instead, a dense ZrO2 layer formed on the ZSZ composite with the highest ZrC content. This indicated that the addition of ZrC was effective for improving the oxidation resistance of ZS composites. 2. Dense ZrO2 layers formed on ZSZ composites only after oxidation above 2000 \\x0e C. Meanwhile, the porosity of the ZrO2 layer increased with decreasing ZrC content in the ZSZ composites. 3. From the volatility diagrams for the oxidation of ZSZ composites, ZrC was concluded to oxidize preferentially over ZrB2 and SiC, even above 2000 \\x0e C. The disappearance of the SiC-depleted layer was caused by the oxidation of ZrC in the SiC-depleted layer, which then became a porous ZrO2 layer. 4. Thermodynamic analysis suggested that the formation of SiO2 is limited during the oxidation of ZSZ composites above 2000 \\x0e C because of the increased O2 partial pressure required to form SiO2 and because the activation energy for the oxidation of ZrC was higher than that for ZrB2 and SiC, unlike the oxidation of ZSZ composites below 2000 \\x0e C. However, because of the longer cooling time relative to the holding time at the maximum temperature, SiO2 is formed during cooling and it remains in oxide layer. 5. The effect of the addition of ZrC to ZrB2eSiC could be evaluated only by oxidation tests above 2000 \\x0e C. The fast heating method proposed in the present study was effective for evaluating the oxidation mechanism of ZSZ composites above 2000 \\x0e C.  References  [4]  [3]  formation  [1] W.G. Fahrenholtz, Thermodynamic analysis of ZrB2-SiC oxidation: of a SiC-depleted region, J. Am. Ceram. Soc. 90 (2007) 143e148. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347e1364. F. Monteverde, R. Savino, M. De Stefano Fumo, Dynamic oxidation of ultrahigh temperature ZrB2-SiC under high enthalpy supersonic flows, Corros. Sci. 53 (2011) 922e929. F. Monteverde, R. Savino, M. De Stefano Fumo, A. Di Maso, Plasma wind tunnel testing of ultra-high temperature ZrB2-SiC composites under hypersonic reentry conditions, J. Eur. Ceram. Soc. 30 (2010) 2313e2321. J. Han, P. Hu, X. Zhang, S. Meng, W. Han, Oxidation-resistant ZrB2-SiC composites at 2200\\x0e C, Compos. Sci. Technol. 68 (2008) 799e806. S.N. Karlsdottir, J.W. Halloran, Rapid oxidation characterization of ultra-high temperature ceramics, J. Am. Ceram. Soc. 90 (2007) 3233e3238. [7] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the  [5]  [6]  [8]  [9]  [15]  [10]  J. Eur.  and HfC,  oxidation of zirconium diborideesilicon carbide in air up to 1500\\x0e C, Ceram. Soc. 27 (2007) 2495e2501. F. Monteverde, R. Savino, Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions, J. Eur. Ceram. Soc. 27 (2007) 4797e4805. E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2and HfB2-based ultra-high temperature ceramics: effect of Ta additions, J. Mater. Sci. 39 (2004) 5969e5977. P. Hu, X.-H. Zhang, J.-C. Han, X.-G. Luo, S.-Y. Du, Effect of various additives on the oxidation behavior of ZrB2-based ultra-high-temperature ceramics at 1800\\x0e C, J. Am. Ceram. Soc. 93 (2010) 345e349. [11] Y. Arai, R. Inoue, H. Tanaka, Y. Kogo, K. Goto, In-situ observation of oxidation behavior in ZrB2-SiC-ZrC ternary composites up to 1500\\x0e C using hightemperature observation system, J. Ceram. Soc. Jpn. 124 (2016) 890e897. [12] Y. Kubota, H. Tanaka, Y. 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Sci. 98 (2015) 233e239. [25] H.-L. Liu, J.-X. Liu, H.-T. Liu, G.-J. Zhang, Changed oxidation behavior of ZrB2SiC ceramics with the addition of ZrC, Ceram. Int. 41 (2015) 8247e8251. [26] M.W. Chase Jr., NIST-JANAF Thermochemical Tables, fourth ed., American Institute of Physics, Woodbury, NY, 1998. J.B. Berkowitz-Mattuck, High-temperature oxidation III. Zirconium and hafnium diborides, J. Electrochem. Soc. 113 (9) (1966) 908e914. [28] A.K. Kuriakose, J.L. Margrave, The oxidation kinetics of zirconium diboride and zirconium carbide at high temperatures, J. Elec. Soc. 111 (1964) 827e831. J.W. Hinze, H.C. Graham, The active oxidation of Si and SiC in the viscous gasflow regime, J. Electrochem. Soc. 123 (1976) 1066e1073. [30] Y. Kubota, H. Hatta, T. Yoshinaka, T. Goto, T. Rong, Use of volume element methods to understand experimental differences in active/passive transitions and active oxidation rates for SiC, J. Am. Ceram. 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},{
  "_id": 111,
  "PDF": "Initial stage of oxidation process and microstructure analysis of HfB2–20 vol._ SiC composite at 1500°C.pdf",
  "Text": "['Available online at www.sciencedirect.com  Scripta Materialia 64 (2011) 617-620  www.elsevier.com/locate/scriptamat  Initial stage of oxidation process and microstructure analysis of HfB2-20 vol.% SiC composite at 1500 °C  De-Wei Ni,a,b Guo-Jun Zhang,a,⇑  Fang-Fang Xua and Wei-Ming Guoa  aState Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics,  Chinese Academy of Sciences, Shanghai 200050, China bGraduate School of  the Chinese Academy of Sciences, Beijing 100049, China  Received 31 October 2010; accepted 3 December 2010  Available online 9 December 2010  The initial stage of oxidation process of HfB2-20 vol.% SiC composite at 1500 °C in air was investigated. With no holding, the oxide scale is composed of a discontinuous SiO2-rich glass layer and an imperfect SiC-depleted layer. Detailed analysis showed that the imperfect SiC-depleted layer contained an HfB2 matrix with partially oxidized HfB2 and SiC particles enclosed in graphite, which revealed that the formation of the SiC-depleted layer during oxidation resulted from the active oxidation of SiC with C as an initial product.  Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: HfB2-SiC; Oxidation; Microstructure; Transmission electron microscope  Ultrahigh-temperature ceramics (UHTCs) have been proposed as candidates for applications such as thermal protection materials on hypersonic aerospace vehicles and reusable atmospheric re-entry vehicles [1,2]. Among the UHTCs, hafnium diboride (HfB2) has a desirable combination of high melting temperature (>3000 °C), excellent resistance to chemical attack and other physical properties, which make it a most suitable candidate for the above applications [3]. In thermal protection applications on hypersonic flight vehicles, UHTCs will be exposed to high temperatures (1500 °C and above) and oxidizing environments [4]. Exposure of HfB2 to air results in the stoichiometric oxidation to HfO2 and B2O3 by reaction (1): HfB2 þ 5 O2 ðgÞ ! HfO2 þ B2O3 ðlÞ 2 Below \\x181100 °C, the B2O3(l) forms a continuous layer, with the oxidation rate appearing to be controlled by the transport of oxygen through the B2O3(l) [5]. Due to its high vapor pressure, the B2O3 evaporates above 1100 °C [6]. The removal of B2O3 by vaporization leaves behind a non-protective porous HfO2 scale, resulting in rapid linear oxidation kinetics above 1400 °C. the [7] The addition of SiC has been reported to improve the oxidation resistance of HfB2 at 1100 °C and above [8].  ð1Þ  ⇑ Corresponding author. E-mail: gjzhang@mail.sic.ac.cn  Above 1100 °C, SiC reacts with O2 to form SiO2 according to the following reaction:  SiC þ 3 2  O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ  ð2Þ  The silica-containing scale on HfB2-SiC is stable at higher temperatures than the B2O3 scale on HfB2 due to its much lower volatility and higher viscosity [9]. Therefore, HfB2-SiC exhibits slow, diffusion-controlled mass gain kinetics over a much greater temperature range than pure HfB2. In recent years, a great deal of research has been conducted on the oxidation behavior of ZrB2-SiC and HfB2-SiC composites [7-15]. Many results indicate that the typical structure of oxide scales on MB2-SiC ceramics after oxidation at \\x181500 °C is (M = Zr, Hf) composed of three layers: (i) an SiO2-rich glassy layer; (ii) a thin MO2-SiO2 layer; and (iii) a SiC-depleted layer [7,9-11,13-15]. Usually, the MO2-SiO2 layer is difficult to detect or separate from the SiC-depleted layer because it is so thin. Layers depleted of SiC have also been observed under the more severe conditions imposed by arc heater testing for both ZrB2-SiC and HfB2-SiC [8,16]. Additionally, Carney et al. [11] showed that the third inner layer of the oxide scale of ZrB2-SiC ceramics 1400-1600 °C was at constituted by a ZrO2 matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions. Fahrenholtz [7] analyzed the development of the layered oxide scale structure using volatility  1359-6462/$ see front matter Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2010.12.002  \\x0c', '618  D.-W. Ni et al. / Scripta Materialia 64 (2011) 617-620  diagrams, which indicated that the formation of the SiCdepleted layer was due to active oxidation of SiC (reaction (3)). Moreover, oxidation studies conducted by Guo and Zhang [15] suggested that the extended SiC-depleted layer (including the thin MO2-SiO2 layer) was most indicative for evaluating the oxidation resistance.  SiC þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  ð3Þ  However, to the best of our knowledge, there has been only limited experimental work involving the development of the oxide scale of HfB2-SiC composites. In particular, no publications have reported the development process of the SiC-depleted layer by experiment. In this study, the initial stage of the oxidation process of HfB2-20 vol.% SiC ceramic in air at 1500 °C was investigated. Emphasis was put on the microstructure evolution of the oxide scale, especially the development process of the SiC-depleted layer. In addition, thermodynamic analysis was used to discuss the observed results. Commercially available a-SiC (D50 = 0.45 lm, 98.5% purity; Changle Xinyuan Carborundum Micropowder Co. Ltd., China) and synthesized HfB2 [17] were used as the starting materials. The starting mixtures (HfB2- 20 vol.% SiC (HS20)) were mixed for 24 h in plastic bottles using absolute ethanol and SiC balls as medium, and dried by rotary evaporation at 70 °C. The mixed powders were then sieved through a 200-mesh screen and placed in a graphite die coated with BN and lined with graphite foil. Subsequently, the composites were hot pressed (HP) at 2000 °C/30 MPa for 1 h with a heating rate of 10 °C min-1. At temperatures below 1600 °C, the furnace was heated under vacuum. Above that temperature, the atmosphere was switched to flowing argon gas. Samples with dimensions of 3 mm \\x02 4 mm \\x02 6 mm were cut from the hot-pressed billets and polished to a 1.5 lm surface finish for oxidation testing. Oxidation studies were conducted in a box furnace by exposing the specimens in stagnant air at 1500 °C without holding. The heating rate was 10 °C min-1. The sample was placed on a zirconia plate with minimal contact area to avoid interaction at high temperature. Microstructures were characterized using scanning electron microscopy (SEM) imaging in an electron probe microanalyzer (JEOL JXA-8100F, Japan) along with energy-dispersive spectroscopy (EDS; Oxford Instruments, UK) and wavelength-dispersive spectroscopy (WDS) for chemical analysis. To better understand the microstructure details, the phase analysis and characterizations were also performed using a 200 kV transmission electron microscope (TEM; JEM2100, JEOL, Japan) along with selected area electron diffraction (SAED) and EDS. The bulk density of the sintered HS20 composite was 9.54 g cm-3, as measured by the Archimedes test method. Using the rule of mixtures and assuming the true densities of HfB2 and SiC to be 11.21 and 3.23 g cm\\x003, the theoretical density was determined to be 9.61 g cm\\x003, giving a theoretical density of >99% for the sample. Figure 1 shows the microstructure of the HS20 ceramic prior to oxidation. The microstructure was typical of those presented previously in the literature. The SiC particles (black phase) were well distributed throughout the HfB2 (gray phase) matrix, and located primarily at HfB2 grain junctions. As a consequence of particles being  pulled out during polishing, a number of pits (similar appearance to pores) were apparent in a scanning electron microscopy (SEM) image. The HfB2 grain size equivalent diameter) was \\x182.6 lm in the as(circular sintered HS20 composite, as determined using an image analysis software package (image-pro 5.0). Figure 2 shows the cross-sectional image and compositional maps of the polished HS20 ceramics after oxida1500 °C. Even without holding 1500 °C, tion at at a discontinuous outer glassy layer was formed, and the oxide scales showed a layered structure. It is widely accepted that the formation of a glassy layer on the surface of MB2-SiC can limit oxygen diffusion to the bulk material and suppress further oxidation. According to the ZrB2-SiC oxidation results by Fahrenholtz [7], when the temperature approaches 1500 °C, the rate of SiC oxidation increases, leading to the formation of SiO2(l). However, the parabolic mass gain kinetics observed in this temperature regime suggest a continuous evolution in the composition of the protective oxide layer from mainly B2O3(l) below 1100 °C to mainly SiO2(l) by 1500 °C rather than the complete loss of B2O3(l) [7]. The detailed WDS analysis in this work also confirmed the existence of B in the outermost layer. Since oxidation is a dynamic process, newly formed B2O3 is expected to dissolve continuously into the glassy layer.  Figure 1. SEM micrograph of a polished HfB2-SiC specimen prepared by hot pressing at 2000 °C for 1 h.  Figure  2. SEM micrograph and compositional maps of a polished section of HS20 composite after oxidation at 1500 °C without holding.  \\x0c', 'D.-W. Ni et al. / Scripta Materialia 64 (2011) 617-620  619  The glassy layer should thus contain some B2O3 at any time during oxidation. In addition, a cluster of HfO2 particles was also observed in this layer. Therefore, the surface glassy layer should be a borosilicate glass, composed of HfO2-SiO2-B2O3. Similar results have also been reported in the ZrB2-SiC system [11,12,15]. Moreover, based on the calculated isothermal section of the 1500 °C, ZrO2-SiO2-B2O3 (BSZ) phase diagram at Guo and Zhang [15] suggested that the ZrO2 first dissolved into the SiO2-B2O3 liquid to form a BSZ liquid at the inner side of the glassy layer. This BSZ liquid would then flow toward the top of the glassy layer. When the B2O3 was lost by evaporation at the outer surface, ZrO2 precipitated from the BSZ liquid. And, as the oxidation progressed, it tended to form rod-shaped ZrO2 inclusions in the glassy layer. According to previous studies, the reactive zone beneath the outer SiO2-rich glass is a so-called SiC-depleted layer, which consists of unoxidized diboride, but may contain some retained SiC [7]. Under the conditions of this experiment, the underlying layer was composed of Hf, Si and O, etc. (Fig. 2). Being subjected to a high temperature for an extremely short time, the discontinuous SiO2-rich glass scale was very thin, and the SiCdepleted layer was imperfectly formed. Consequently, it was difficult to separate these layers from each other. However, this microstructure might provide some useful information for the initial development of the SiCdepleted layer and oxide scale. In Figure 2 the red line shows the rough interface separating the oxide layers from the unaffected bulk. The whole thickness of the oxide layers was approximately 10 lm. As oxidation proceeded, the gaseous products produced by the above-mentioned reactions near the interface would accumulate continuously. As their pressure exceeded the ambient pressure, these gases would move toward the outside, causing the thickness of the outer borosilicate glass and underlying SiC-depleted region to increase continuously. At the same time, the interface (the red line shown in Fig. 2) would move inward. To better understand the development of oxide scale, especially the formation of the SiC-depleted layer, the microstructure of the above-oxidized specimen was further analyzed with a TEM. As the oxide scale in this experiment was thin, the TEM view was just of the innermost reaction zone. Figure 3a shows a bright field image of the typical microstructure of the SiC-depleted region. In the figure, the grayish-black phase is HfB2 and the white phase with a stripe is SiC. The SiC particles exhibited a unique core-shell structure, being surrounded by a semitransparent glass-like material. EDS analysis (Fig. 3b) indicated that this semitransparent glass-like material was composed of carbon. In agreement with the EDS, SAED (inset) further revealed that it was polycrystalline graphite. However, there were also weak signals of O, “Hf and S” due to disturbance from the adjacent grains. Obviously the SiC grains were partially oxidized, graphite being a reaction product of SiC oxidation. The distribution of elements in the above region, including the partially oxidized SiC, is shown in Figure 4. Unlike Figures 3a and 4 is a scanning transmission electron microscope (STEM) image, which had an inverse contrast. The graphite layer surrounding the SiC parti cle could be identified easily from the composition maps. A weak oxygen signal was also seen to exist at the position of the HfB2 particles, revealing that the HfB2 was partially oxidized. EDS point analysis on HfB2 particles further confirmed this (not shown here). The above feature of the innermost oxide zone (i.e. the SiC-depleted layer) was slightly different from previous results [7- 11,13-16]; this difference is accounted for below. According to previous results and thermodynamic analysis, the oxygen activity (oxygen partial pressure, P O2 ) beneath the SiO2-rich scale should be significantly ðP O2 \\x18 104Pa in airÞ less than the ambient atmosphere in air) because of the chemical potential gradient associated with any diffusion profile. It is generally acknowledged that the low oxygen partial pressure results in the active oxidation of SiC, leading to the formation of an SiC-depleted layer (composed of unoxidized diborides and a small amount of SiC). Unlike in the above-mentioned analysis and previous results, solid C (polycrystalline graphite) was observed in the SiCdepleted layer, which is reported for the first time. Others have widely reported an SiC-depleted region just above the unoxidized MB2-SiC bulk in samples heated between 1500 and 2200 °C [7-11,13-16], while only Monteverde [18] and Carney et al. [11] showed C-rich inclusions in the same layer. We know that previous analysis has shown that low oxygen partial pressure beneath an SiO2-rich scale results in active oxidation of SiC by reaction (3). Combining this with our own findings, we propose another possible active oxidation reaction of SiC: SiC þ 0:5O2 ðgÞ ! SiOðgÞ þ C  ð4Þ  Thermodynamic calculation indicated that reaction (4) was favorable across the whole temperature range studied. However, carbon was only thermodynamically stable when the oxygen partial pressure was below 5 \\x02 10\\x0015 Pa (which was much lower than the actual oxygen partial pressure in this reaction region) [7,19]. Therefore, the only rational explanation was that the oxidation of carbon (reaction (5)) was much slower than reaction (4). Also, the short oxidation duration resulted in the existence of carbon in the SiC-depleted layer. With prolongation of the holding time at high temperature, the residual carbon would be removed by oxidation to CO, and a porous SiC-depleted layer would be formed:  C þ 0:5O2 ðgÞ ! COðgÞ  ð5Þ  On the other hand, previous results have widely recognized that the oxygen activity in the SiC-depleted layer is low enough to prevent oxidation of MB2 at 1500 °C. Therefore, the partially oxidized HfB2 in the inner reaction region should be oxidized before the protective SiO2-rich layer is formed. However, as stated above, the oxidation is a dynamic process. Even after a continuous SiO2-rich layer is formed, the partially oxidized HfB2 near the inner side of the glassy layer would further oxidize as oxidation progresses. Studies conducted at different temperatures or with certain time intervals might help us further understand the reaction mechanisms during oxide scale formation. In summary, the oxidation behavior of a dense HfB2- 20 vol.% SiC composite produced by hot pressing at  \\x0c', '620  D.-W. Ni et al. / Scripta Materialia 64 (2011) 617-620  Figure 3.  (a) TEM micrograph of the innermost oxide zone containing partially oxidized SiC. The SiC particles exhibited a core-shell structure, being  surrounded by well-crystallized graphite. The SAED patterns of the SiC and graphite are inset. (b) EDS analysis of graphite and SiC.  Figure 4. STEM image and elements distribution (Si, C, Hf, O and B) of the innermost oxide zone containing partially oxidized SiC.  2000 °C using synthesized HfB2 powder was investigated. After oxidation in air at 1500 °C without holding, an discontinuous outer SiO2-rich glass layer and an imperfect underlying SiC-depleted layer were formed that were difficult to separate from one other. Detailed TEM analysis showed that SiC particles in the SiC-depleted layer exhibited a unique core-shell structure surrounded by polycrystalline graphite, which revealed that active oxidation of SiC in the SiC-depleted layer first involved the reduction to C. This work has provided experimental evidence for the development of an SiC-depleted layer.  Financial support from the Chinese Academy of Sciences under the Program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program), The National Natural Science Foundation of China (No. 50632070), The Science and Technology Commission of Shanghai (Nos. 08520707800 09ZR1435500) and the CAS Special Grant for Postgraduate Research, Innovation and Practice is greatly appreciated. The authors are grateful to Mei-Ling Ruan and Xiang Gao for help with the TEM analysis.  [1] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, J. Mater. Sci. 39 (2004) 5887. [2] F. Monteverde, A. Bellosi, S. Guicciardi, J. Eur. Ceram. Soc. 22 (2002) 278. [3] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, J. Am. Ceram. Soc. 90 (2007) 1347.  J. Am.  J.  Eur.  J. Am. Ceram.  [4] A. Bongiorno, C.J. Fo¨ rst, R.K. Kalia, J. Li, J. Marschall, A. Nakano, M.M. Opeka, I.G. Talmy, P. Vashishta, S. Yip, MRS Bull. 31 (2006) 410. [5] J.B. Berkowitz-Mattuck, J. Electrochem. Soc. 113 (1966) 908. [6] J. Li, T.J. Lenosky, C.J. Fo¨ rst, S. Yip, Soc. 91 (2008) 1475. [7] W.G. Fahrenholtz, J. Am. Ceram. Soc. 90 (2007) 143. [8] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, J. Mater. Sci. 39 (2004) 5925. [9] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Ceram. Soc. 89 (2006) 3240. [10] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Ceram. Soc. 27 (2007) 2495. [11] C.M. Carney, P. Mogilvesky, T.A. Parthasarathy, J. Am. Ceram. Soc. 92 (2009) 2046. [12] S.N. Karlsdottir, J.W. Halloran, J. Am. Ceram. Soc. 92 (2009) 481. [13] D.W. Ni, G.J. Zhang, Y.M. Kan, Y. Mater. 60 (2009) 615. [14] D.W. Ni, G.J. Zhang, Y.M. Kan, Y. Mater. 60 (2009) 913. [15] W.M. Guo, G.J. Zhang, 2387. [16] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D. Ellerby, Refractory Appl. Trans. 1 (2005) 1. [17] D.W. Ni, G.J. Zhang, Y.M. Kan, P.L. Wang, Ceram. Soc. 91 (2008) 2709. [18] F. Monteverde, App. Phys. A 82 (2006) 329. [19] A.H. Heuer, V.L.K. Lou, J. Am. Ceram. Soc. 73 (1990) 2785.  J. Eur. Ceram. Soc. 30 (2010)  Sakka,  Scripta  Sakka,  Scripta  J. Am.  \\x0c']"
},{
  "_id": 112,
  "PDF": "Interfacial reaction on oxidation of carbides with formation of carbon.pdf",
  "Text": "['Solid State Ionics 141 - 142  Ž  2001  .  99 - 104  www.elsevier.comrlocaterssi  Interfacial reaction on oxidation of carbides with formation of carbon  Shiro Shimada)  DiÕision of Materials and Engineering, Graduate School of Engineering, Hokkaido UniÕersity, West-8, North-13, Kita,  Sapporo 060-8628, Japan  Abstract  Ž  .  \\x04  4  \\x04  4  .  Ž  P  O 2  The oxidation of carbides ZrC, HfC, TiC using single crystals with the 200 or 220 face was carried out at temperatures of 500 - 1500 8C and oxygen partial pressures of 0.02 - 80 kPa. The oxidation proceeded linearly with time. The orientation of crtor m-ZrO , m-HfO preferred 200 or 220 and Ti O was recognized on the oxidized surfaces of the crystals or near the reaction interface. The oxide scale consisted of inner dense dark scale zone 1 and outer porous whitergray scale zone 2 . Zone 1 contained much carbon at around 23 - 25 at.%. Simultaneous TG -DTA -MS analysis of HfC powders showed that the oxidation of Hf component in HfC occurs to produce carbon at 380 - 600 8C above which the ZrCrzone 1 interface, remaining carbon is oxidized with evolution of CO . From HRTEM observation at it was found that c-ZrO crystallites are formed directly on the ZrC lattice while maintaining the above preferred orientation. The interfacial reaction with the formation of carbon on oxidation of the carbides was discussed from the above results. q 2001 Elsevier Science B.V. All rights reserved.  Ž  .  .  3  5  Ž  2  2  2  2  Keywords: Oxidation; Carbon; Kinetics; Mechanism; Interfacial reaction; Carbide; Preferred orientation; Compounds, ZrC, HfC, TiC, ZrO ,2 HfO , TiO , Ti O , ZrC O , HfC O  1y x  x  1y x  x  2  2  3  5  1. Introduction  Ž  There have been many studies concerning oxidation of transition metal carbides TiC, ZrC, HfC using powder and sintered pellet samples. Of these, some authors including Shimada suggested the formation of carbon during oxidation of ZrC, HfC and NbC powders 1 - 4 . The formation of carbon was inferred from exceeding 100% of a degree of oxidation, which was obtained by dividing the measured  .  w  x  )  Fax: q81-11-706-6576. E-mail address: shimashi@eng.hokudai.ac.jp  Ž  S. Shimada .  .  weight gain by the theoretical one, which was calculated by assuming the complete conversion of MeC to MeO with formation of CO according to Eq. .1 :  Ž  2  2  MeC q O s MeO q CO  2  2  2  1Ž  .  w  x  where Me s Zr, Hf. For example, Shimada et al. 2 reported that the oxidation of HfC powders yields a degree of 123% at 5378C, then decreases slowly and attains final 100% oxidation at extended times. The author studied the oxidation of ZrC and HfC single crystals at 500 - 1500 8C and at various oxygen pressures, and examined the oxide scale by backscattering electron images BEI in a scanning electron  Ž  .  0167-2738r01r$ see front matter q 2001 Elsevier Science B.V. All PII: S 0 1 6 7 2 7 3 8 0 1 0 0 7 2 7 5  Ž  .  rights reserved.  \\x0c', '100  S. Shimada r Solid State Ionics 141 - 142  2001 (  ) 99 - 104  Ž  .  x  w  x  w  .  Ž  microscope SEM , showing that the oxide scale consists of two sub-scales with the inner and outer scales growing parabolically and linearly, respectively 5,6 . He also observed the interface between the ZrC and zone 1 by high-resolution transmission electron microscope HRTEM , evidencing that zone 1 contains a large amount of elemental carbon 7 . However, he did not present any detailed information on what interfacial reaction occurs to produce carbon during oxidation of the carbides, nor gave any direct evidence on oxidation of the metal component in the carbide prior to that of the carbon component. This paper describes the interfacial reaction occurring during oxidation of the single crystal carbides of ZrC, HfC and TiC with formation of carbon without oxidation. The evolution of CO gas during oxidation of the HfC powder sample was followed in order to show the oxidation of the carbon remaining unoxidized after the oxidation of Hf component in HfC.  2  2. Experimental  Ž  Ž  .  Single crystals and powders of the carbide ZrC, HfC, TiC were used for the oxidation studies. The crystals did not contain any free carbon. The single crystals were cleaved into several plates parallel to the 200 or 220 plane and then polished mechanically with 3 mm diamond paste. A crystal of the carbide was put in an alumina crucible and oxidized isothermally at temperatures of 500 - 1500 8C at oxygen pressures of 0.02 - 80 kPa. The weight gain of the crystal was monitored using an elecro-microbalance SETARAM, Model TG92 . The phases formed on surfaces of oxidized crystals were identified by X-ray diffraction XRD . The oxidized crystals were cut and polished for cross-sectional observation by BEI-SEM Jeol JSM 35 CF . Wavelength-dispersive WDS X-ray microanalysis XMA Jeol JCMA 733 was conducted to quantitatively determine the atomic concentrations of Zr, Hf, or Ti, O and C in the oxide scale. A thin foil of a ZrC crystal oxidized at 6008C was prepared for TEM observation operating at 300 kV JEOL, JEM-3010 . Chemical analyses for the O, C, and Zr, Hf, or Ti elements were performed on the thin foil by EDX.  Ž .  . Ž  Ž  .  Ž  .  Ž  .  Ž  .  .  The powder oxidation of HfC was followed by simultaneous measurements of the weight and thermal changes and the CO evolution using TGA - DTA -MS TGA -DTA -MS 2000, MAC Science; VG Gas Analysis, Fisons Instrument at a heating rate of 108C miny1 and at an oxygen pressure of 10 kPa. The sample weight was about 10 mg.  Ž  .  2  3. Results and discussion  DWrW  Ž  .  o  The weight gain of a ZrC crystal was monitored during isothermal oxidation at 5008C, 5508C and 6008C for times up to 240 h 5 . An almost linear weight gain was observed, but it gradually slowed down at extended times. The oxidation above 9008C occurred linearly. It was found that the oxidation of a ZrC crystal is described by a linear kinetic equation Eq. 2 :  ..  Ž  Ž  w  x  DWrW s k  o  t  L  2Ž  .  L  t  k  is  the  and  rate  constant  the oxidation  where time. The weight gains by oxidation of an HfC crystal at 600 - 900 8C were followed. Oxidation occurred very slowly at 6008C and proceeded much faster at temperatures higher than 7008C. The weight gains at 7758C, 8508C, and 9008C exceeded 10.5 wt.% calculated based on Eq. 1 , which corresponds to 100% oxidation. The high temperature oxidation above 9008C also proceeded linearly. Oxidation of a TiC crystal at temperatures higher than 7008C was obeyed by a similar linear kinetic. It is concluded that the oxidation of the single crystals of TiC, ZrC and HfC proceeds linearly with time. Exceeding 100% oxidation for oxidation of a HfC crystal is due to the retention of carbon according to Eq. 3 :  Ž  .  Ž  .  HfC q O s HfO q C  2  2  3Ž  .  Figs. 1 and 2 show the XRD patterns on the oxidized crystal surface of ZrC and HfC, respecP s 80 - 0.02 tively, at 700 - 1500 8C and at kPa. The mixed phases of tetragonaland monoclinic-ZrO2 tand m-ZrO appear, with weak orientation of 110 or 200 peak of t-ZrO , independent of the oxygen pressure at 7008C Fig. 1 a - c . The mZrO peaks increase their intensities with increasing  ..  O 2  Ž  Ž  .  Ž  Ž  .  2  2  2  \\x0c', \"S. Shimada r Solid State Ionics 141 - 142  2001 (  ) 99 - 104  101  Fig. 1. XRD patterns of the oxidized crystal surface of ZrC at 700 - 1500 8C. 7008C at 15008C; B, t-ZrO , ', m-ZrO ;  0.02 kPa; P s 0.02 kPa at the figure in parentheses corresponds to the intensity referred by the JSPDS card.  13008C,  11008C,  0.2 and  Ž .  2,  O 2  b  d  a  c  e  Ž  .  Ž  .  Ž  .  Ž  .  Ž  .  f  2  2  Ž  Ž  .  Ž ..  2  2  Ž  temperature from 11008C to 15008C Fig. 1 d - f . It is thus recognized that the preferred 200 or 220 trm-ZrO orientation of occurs on a ZrC crystal oxidized at 700 - 1500 8C. The 200 line of monoclinic-HfO m-HfO is seen with a shoulder of 002 or 020 at 7008C and 9008C at 80 kPa Fig. 2 b , c , and at 9008C at 0.08 kPa Fig. 2 e , but no peak except for the 200 HfC appears at 7008C and 0.08 kPa Fig. 2 d . The 002 line of m-HfO becomes more intensified in spite of little growth of the other lines above 11008C Fig. 2 f - h . Oxidation of a  Ž ..  Ž .  . Ž  ..  ..  ..  Ž  Ž  .  Ž  Ž  Ž  Ž  Ž  2  2  2  .  Ž  TiC crystal gave the 110 and 220 diffraction lines of TiO rutile at 7008C, which were not greatly changed at 900 - 1500 8C. A TiC crystal oxidized at 15008C was polished thinly at a regular interval of about 5 mm from the outermost surface towards the interface. The XRD patterns for this polished crystal were largely unchanged down to a depth of 40 mm, but a very small 020 diffraction line of Ti O appeared at a depth of 45 - 50 mm. This line greatly increased towards the interface, in contrast to the 110 line of TiO , which began to decrease at about 60  3  5  2  \\x0c\", '102  S. Shimada r Solid State Ionics 141 - 142  2001 (  ) 99 - 104  Ž  .  x  w  Ž  Ž  .  Ž  Ž  Ž  .  Ž  .  Ž  .  Ž  .  Ž  ..  ..  with dark contrast zone 1 and an outer gray scale zone 2 . A feature of the two scales is that zone 1 has dense structure with few pores and zone 2 is porous and cracked. Thin, wavy lamella scale zone 1 and porous, cracked scale zone 2 are seen in the oxide scale of TiC. The lamella structure consists of black and gray alternate strips parallel to the interface. In previous papers 8,9 , it was reported that zone 1 contains as much as 23 - 25 at.% carbon and zone 2 contains 4 - 11 at.% carbon and that elemental carbon is deposited in zone 1 during oxidation of the carbide. Non-isothermal oxidation of the HfC powder was 108C miny1 followed at a heating rate of at an oxygen pressure of 10 kPa by simultaneous TGA - DTA -MS analysis Fig. 4 . From TGA result Fig. 4 A , the weight gain begins at about 3808C, continuing with rising temperature, going up to a maximum of 160% at 6708C through 100% oxidation, and finally returns to 100% at 7608C. One hundred sixty percent oxidation means no consumption of carbon according to Eq. 3 . DTA curve Fig. 4 B shows that oxidation begins exothermally at 3808C, making two overlapping exothermic peaks a low and high temperature DTA peaks at 6108C and 6808C, and finishes at 7608C. From MS analysis Fig. 4 C , CO begins to evolve at 5708C, corresponding to 60% oxidation on the TG curve, and finishes at 7608C. The CO evolution peak corresponds to the high temperature DTA peak. It is thus suggested that the low temperature DTA peak is due to the oxidation of sole Hf in HfC, forming HfO and carbon, and the high temperature one due to the subsequent oxidation of carbon with the CO evolution. It should be stated that the preceding oxidation of Zr in ZrC with subsequent oxidation of carbon was similarly observed for ZrC powders by TG -DTA -MS analysis. The interfacial area between ZrC and zone 1 was observed by HRTEM Fig. 5 , in order to follow the interfacial reaction with formation of carbon. As reported by a previous paper 7 , it is seen that amorphous carbon exists along the interface, with many c-ZrO crystallites of 4 - 10 nm sizes distributed in a regular way. It is interesting to see that some c-ZrO crystallites are bonded directly to the ZrC lattice see two arrows , with the direction of their lattice fringes parallel to that of the ZrC lattice  ..  Ž  .  Ž  .  Ž  Ž  Ž  .  w  x  2  2  2  2  2  2  Fig. 2. XRD patterns of the oxidized crystal a Unoxidized; P s 80 kPa at 700 - 1500 8C. 9008C; P s 0.08 kPa at d 7008C, e 9008C, 15008C; ^, HfC. 13008C, h  O 2  O 2  Ž  .  Ž  .  Ž  Ž  .  Ž  .  b  surface of HfC at  . Ž .  f  7008C and  11008C,  Ž Ž  c  g  . .  5  3  4  \\x04  4  \\x04  mm depth, and disappeared at a depth of 105 mm. At this depth, only Ti O with a very intensified 020 line was observed. It is concluded that when the 200 or 220 face of the carbides is oxidized, the preferred 200 or 220 orientation of the oxide scale trmZrO , m-HfO , Ti O occurs. Cross-section of the scale obtained by oxidation of the carbides at 600 - 1500 8C is depicted in Fig. 3 6 - 8 in which three layers are recognized in all the oxide scale. A white contrast represents the carbide, on which two sub-scales are formed; an inner scale  Ž  .  w  x  2  2  3  5  \\x0c', 'S. Shimada r Solid State Ionics 141 - 142  2001 (  ) 99 - 104  103  Fig. 3. BEI-SEM images of cross-sections of ZrC, HfC and TiC crystals oxidized at 6008C, 7008C and 15008C, respectively. A ZrC,  Ž  .  Ž  .  B  HfC,  Ž  .  C  TiC.  Ž  .  Ž  .  Ž  .  Ž  .  2  2  4  \\x04  4  \\x04  such as 200 and 220 ZrC or 200 and 220 c-ZrO . The preferred 200 or 220 orientation of trm-ZrO , m-HfO , and Ti O on the oxide scale is explained by the fact that c-ZrO crystallites are formed while maintaining the crystallographic relationship between c-ZrO and the ZrC. The relative concentration profile of oxygen across the interface  2  3  5  2  2  Fig. 4. TG -DTA -MS analysis for oxidation of HfC powder. A C MS curve; Heating rates108C TG curve, B DTA curve, miny1 , P s10 kPa, the sample weights10 mg.  Ž  .  Ž  .  Ž  .  O 2  over the distance of 300 nm determined by EDX analysis is depicted in Fig. 6. The oxygen content in zone 1 decreases from a distance of 100 nm away from the interface, indicative of occurring oxygen diffusion, assuming that zone 1 acts as a barrier for the oxygen diffusion. A relatively small amount of oxygen is also detected in the ZrC lattice, decreasing down to a distance of 200 nm from the interface, implying that oxygen is dissolved in the ZrC, forming oxycarbide, ZrC O , in which decreases with increasing distance from the interface. The formation of HfC O in HfC crystal was also confirmed 8 . Referring to a previous paper 5 , the diffusion of oxygen limits the growth rate of zone 1.  1y x  1y x  x  w  x  w  x  x  x  Fig. 5. HRTEM image at the ZrCrZrO  interface.  2  \\x0c', '2001 (  ) 99 - 104  S. Shimada r Solid State Ionics 141 - 142  104  Ž  .  Fig.  6. The  relative  count  ratio  OrHf  across  the ZrCrZrO2  interface.  It is thus suggested that a local thermodynamic equilibrium is established at the interface, at which a partial oxygen pressure was calculated to be ; 10y2 0 atm; this value can oxidize Me in MeC, but not carbon.  4. Summary: carbides  interfacial reaction on oxidation of  From the above results and discussion, the interfacial reactions between the carbide and zone 1 and between zones 1 and 2 are summarized as follows. The carbiderO interface in early stage. First, oxygen is dissolved in the carbide, forming oxycarbide, MeC O , according to Eq. 4 ; MeC q 3 xr2O s MeC O q x CO  Ž .i  2  Ž  .  1y x  x  4Ž  .  2  1y x  x  2  where x 0.4. carbiderzone .ii The 1 interface. As is increased with the dissolved amount of O in the oxycarbide, the oxycarbide is saturated with oxygen at x ; 0.4, which produces MeO or Ti O crystallites Eq. 5 : O q 1 y xr2 O s MeO q 1 y x C .  Ž  x  2  2  3  5  Ž  Ž  ..  MeC  Ž  .  Ž  .  1y x  x  2  2  5Ž  .  This is the initiation of zone 1. As zone 1 thickens parabolically, it acts as a barrier for the diffusion of oxygen, and the local oxygen activity becomes so low that the carbon of oxycarbide cannot be oxidized but Me is oxidized, leaving elemental carbon at the interface. When zone 1 reaches a critical thickness 2 - 4 mm 5 , it begins to crack due to stress by a difference of thermal expansion coefficient between the carbide and zone 1. This is an initiation of cracked zone 2. Repeating process of the formationr cracking of zone 1 maintains its thickness and thickens zone 2 linearly. This is why the overall weight gain of the carbides by oxidation occurs linearly with time. 1r2 iii The zone interface. Since zone 2 is cracked, O gas can easily penetrate to the front of zone 1 where the remaining carbon is oxidized to evolve CO Eq. 6 :  Ž  w  x.  Ž  .  2  Ž  Ž  ..  2  C zone 1 q O s CO .  6  Ž  .  Ž  .  2  2  CO out  gas makes many pores in zone 2 when it goes through zone 2.  2  References  w x1  L. Dufour, J. Simon, P. Barret, C. R. Acad. Sci., Ser. C 265  Ž  .  1967  171.  w x2  S. Shimada, M.  Inagaki, K. Matsui,  J. Am. Ceram. Soc. 75  Ž  .  1992  2671.  w x3  D. Gozzi, M. Montozzi, P.L. Cignini, Solid State Ionics 123  Ž  .  1999  1.  w x4  D. Gozzi, M. Montozzi, P.L. Cignini, Solid State Ionics 123  Ž  .  1999  11.  w x5  S. Shimada, M. Nishisako, M.  Inagaki, K. Yamamoto, J. Am.  Ž  .  Ceram. Soc. 78  1995  41.  w x6  S. Shimada, K. Nakajima, M. Inagaki, J. Am. Ceram. Soc. 80  Ž  .  1997  1749.  w  x  Ž  .  7  S. Shimada, M.  Inagaki, M. Suzuki, J. Mater. Res. 11  1996  2594.  w x8  S. Shimada, F. Yunazar, S. Otani,  J. Am. Ceram. Soc.  83  Ž  .  2000  721.  w  x  Ž  .  9  S. Shimada, Science Forum accepted .  \\x0c']"
},{
  "_id": 113,
  "PDF": "Investigation of the effects of temperature and oxygen partial pressure on oxidation of zirconium carbide using different kinetics models.pdf",
  "Text": "['Journal of Alloys and Compounds 509 (2011) 2395-2400  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j a l l c o m  Investigation of the effects of temperature and oxygen partial pressure on oxidation of zirconium carbide using different kinetics models  Xin-Mei Hou a,b , Kuo-Chih Chou a,c,∗  a Department of Physical Chemistry, University of Science and Technology Beijing, Beijing 100083, China b The Key Laboratory of Ecological and Recycle Metallurgy, Ministry of Education, University of Science and Technology Beijing, Beijing 100083, China c Department of Materials Sciences and Engineering, Shanghai University, Shanghai 20072, China  a r t i c l e  i n f o  a b s t r a c t  The oxidation kinetics of ZrC materials is an important physicochenmical property for their practical application. Although the oxidation data have been extensively measured, the quantitative relationship of oxidation curves to various factors such as temperature and oxygen pressure is still limited. In this article, the oxidation kinetics of ZrC materials under the conditions of different rate-controlling steps existing has been investigated using two kinds of models (the model used in the literatures and Chou’s model) based on the experimental data available in the literatures. The calculated results show that both models can fit the oxidation data well. Compared with the previous models used in the literatures, Chou’s model can give a clear physical meaning in expressing all parameters. The most important thing is that Chou’s model can give a more accurate performance in the theoretical analysis. © 2010 Elsevier B.V. All rights reserved.  Article history:  Received 6 June 2010 Received in revised form 3 November 2010 Accepted 4 November 2010 Available online 13 November 2010  Keywords:  Zirconium carbide Oxidation Kinetic model Temperature Oxygen partial pressure  1.  Introduction  Most of the transition metal carbides are refractory materials having good mechanical properties particularly useful for cutting tools and drilling heads. For example, zirconium carbide, ZrC is one of the most refractory materials with high melting point (3693 ± 20 K) and good mechanical properties especially a hardness that is comparable with the hardness of TiC [1]. Therefore, ZrC is an important high-temperature structural material and can be applied as a promising nuclear material [2-5]. By comparison, its chemical stability in air and at high temperature is quite poor, which restricts its range of applications. A systematical investigation of the oxidation kinetics for ZrC is essential. Several studies on the oxidation of ZrC both as powder [6-9] and single crystal samples [1,5,10-13] have been reported. Various authors have identified that the oxidation of ZrC is affected by many factors such as temperature, oxygen partial pressure and phase of oxide product. Among these factors, the morphological development of oxide phase is a very important one because it changes the oxidation behavior of ZrC. Therefore, most current works have been focused on the reaction process based on mor ∗ Corresponding author at: Department of Physical Chemistry, University of Science and Technology Beijing, Beijing 100083, China. Tel.: +86 010 6233 2646; fax: +86 010 6233 2570. E-mail addresses: kcc126@126.com, houxinmei@ustb.edu.cn (K.-C. Chou).  0925-8388/$ - see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2010.11.028  phological development and the results showed a complex picture. Bellucci et al. [1] investigated the oxidation behavior of three ZrC single crystals in Ar/O2 mixtures using Auger Electron Spectroscopy (AES), Electron Probe Micro-Analysis (EPMA), Scanning Electron Microscopy (SEM), X-ray Diffraction (XRD) and Raman techniques. The results identified carbon in amorphous state retention during the oxidation process and the presence of both the c-ZrO2 and mZrO2 phases. Shimada et al. [12,13] also confirmed the experimental fact during the oxidation of ZrC single crystal. Besides these, they carried out the oxidation experiments on ZrC powder by simultaneous thermogravimetry (TG), differential thermal analysis (DTA) and mass spectrometry (MS) at vious oxygen pressure between 0.5 and 40 kPa in the temperature range of 293-1273 K. It was pointed out that the oxidation of ZrC overshoots a degree of oxidation of 100% depending on temperature and oxygen pressure, which was due to the formation of carbon during oxidation. For the application of ZrC material, it is important to foresee what damage is caused by high-temperature oxidation as well as to prevent it. Compared with the knowledge of the sequence of events that occur during the oxidation, theoretical investigation especially in quantitative aspect is far more enough. The current models used in the literatures such as the well known Jander’s equation, the parabolic rate law and linear rate law. are usually expressed with a pre-exponential term k and an apparent enthalpy. However, the constant k does not have clear physical meaning and thus these expressions are only empirical fits and used for interpolating the available data. In addition, the current models used in literature are  \\x0c', '2396  X.-M. Hou, K.-C. Chou / Journal of Alloys and Compounds 509 (2011) 2395-2400  Nomenclature  (cid:2)  D0ˇ 0 K 0ˇ 0  vm  \\x01Ed \\x01Er \\x01  PO2 P eq O2  R0 L0 k0 R t T k0  \\x01  ((cid:5)i )mea ((cid:5)i )cal  N  reacted fraction of oxidation  constant independent of temperature  constants independent of temperature coefficient related to the density of reactant and product the apparent activation energy of diffusion the apparent activation energy of chemical reaction temperature-increasing rate oxygen partial pressure in gas phase  oxygen partial pressure in equilibrium with oxide  radius of the original whole particle the original thickness of the pellet a temperature-independent constant gas constant time in second absolute temperature with K temperature-independent constant the average relative error experimental data the value calculated from the model the number of experimental points  established on the following assumption, i.e., d(cid:2)/dt = k(T)f((cid:2)) where f((cid:2)) is a function of reacted fraction (cid:2) (some researcher called it as the reaction model) and k(T) a function of temperature T, normally it can be expressed as an Arrhenius equation, k(T) = A exp( − \\x01E/RT). Apparently, the form of k(T)/f((cid:2)) means that the variables (cid:2) and T can be expressed separately in an expression, which actually is an important assumption for this kind of models. Therefore, the applicability of these models will depend on the reliability of this assumption. In general, this assumption might be acceptable. However big errors will be introduced under the conditions that the reaction is controlled by multiple steps or the sample shape is other than a sphere as pointed out in our previous paper [14]. Recently, Chou’s model [14,15] has been used for the theoretical treatment, which is different from all current models used in the literatures because it is established on a real physical picture with certain size or shape of solid materials undergoing a real process of molecular and atom movement. Since these formulae express the reacted fraction as a function of time, temperature, sample shape, oxygen pressure, etc. in an explicit way, one can easily carry out a calculation and give a good theoretical discussion for practical problem, which has been successfully applied to treat the oxidation behavior of many non-oxide material systems, such as SiAlON and Si3N4 . [14-16]. The aim of this paper is to prevent the oxidation behavior of ZrC material from a quantitative standpoint based on the reported data. The influence factors, especially temperature and oxygen pressure on the oxidation kinetics will be quantitatively investigated for the first time. The results obtained by our new model will be discussed and compared with those reported in the literatures, in terms of which to obtain some idea for the future application and development of ZrC materials.  2.  Experimental  2.1.  Oxidation of ZrC powder  Shimada and Ishii [7] have systematically investigated the oxidation of ZrC powder with the average particle size of 3.3 \\u242em under both non-isothermal and isothermal conditions using an electro-microbalance. The isothermal oxidations were carried out at the temperature range of 653-823 K with different oxygen pressures and the results are shown in Fig. 1a-c. The oxide products were analysed by  XRD and SEM. It showed that at lower temperature, i.e., below 743 K, the oxidation product is amorphous and above 743 K, c-ZrO2 crystalline appeared and the resulting volume expansion growth stress cracks the grains, which led to the oxidation behavior change (Fig. 1b and c). The oxidation kinetics was also investigated by Shimada and Ishii [7], who pointed out that the oxidation was controlled by diffusion and can be well described by the Jander’s equation.  [1 − (1 − (cid:2))1/3 ]  2 = kt  (1)  where (cid:2) is the reacted fraction at time t and k is the rate constant that depends on temperature and pressure. Since at about 743 K, the oxidation mechanism changed because of the phase transition. Eq. (1) was well applied during the whole temperature range with different apparent activation energies, i.e., 138 kJ/mol below 743 K and 180 kJ/mol above that temperature. Fig. 2a and b shows the effect of oxygen pressure on oxidation behavior under both non-isothermal and isothermal conditions. In view of the non-isothermal oxidation behavior of ZrC powder at different oxygen pressures, the oxidation rate increased with oxygen pressure increasing. While at all oxygen pressure, the oxidation reaction began at about 573 K and the rate increased from 700 K (shown in Fig. 2a). Fig. 2b shows the oxidation behavior at 753 K with different oxygen pressures, in which the oxidation rate was faster at higher oxygen pressure and was controlled by diffusion.  2.2.  Oxidation of ZrC pellet  In practice, ZrC material is usually used in the form of compact and the application temperature is usually above 850 K. Kuriakose and Margrave [11] investigated the oxidation of ZrC pellet with the purity of 99.5% in the temperature range of 827-925 K (as shown in Fig. 3). The oxygen partial pressure used in the experiments was 9.8 × 104 Pa. The result showed that the oxidation followed a linear rate law with the activation energy of 70.1 kJ/mol.  3. Results and discussion  From the above experimental results, the oxidation of ZrC material is a complicate process and affected by many factors, such as temperature, oxygen pressure and the shape of ZrC material. A general mechanism of oxidation reaction for ZrC material can be described by a series of steps, i.e., (1) oxygen in the bulk gas phase transfer to the surface of ZrC material; (2) oxygen diffusion through the boundary layer between gas phase and solid phase; (3) physisorption of oxygen molecules; (4) dissociation of oxygen molecules and chemisorption; (5) surface penetration of oxygen atoms; (6) diffusion of oxygen through the oxide product layer to the oxide/ZrC interface; (7) nucleus formation and chemical reaction producing oxide product and gas; (8) gas diffusion through oxide product to the surface of ZrC material; and (9) gas diffusion through the gas/ZrC material boundary to the gas bulk. The nine steps are basically enough to describe the oxidation process of ZrC material. For the purpose of simplification and practical convenience, the rate of limited-step method is usually applied such as Jander’s equation, the linear rate law and Chou’s model. In this section, the experimental data offered by Shimada and Ishii [7] and Kuriakose and Margrave [11] were used to investigate the oxidation kinetics of ZrC materials in terms of the models used in the literatures and Chou’s model, respectively. The obtained results will be compared.  3.1. Application of the models used in the literature  3.1.1. Oxidation kinetics of ZrC powder  In case of the system of ZrC powder offered by Shimada and Ishii [7], Jander’s equation is employed to treat its oxidation kinetics. The value of k can be calculated from the slope line for each temperature. The activation energy is obtained by linear regression of the plot of ln k versus of 1/T, which was calculated to be 142 kJ/mol below 743 K and 175 kJ/mol above that temperature. The different values implied the change of oxidation kinetics caused by the phase transition. In addition, the values of the activation energy obtained are consistent with the results reported in the literatures [7], implying the reasonable of Jander’ equation.  \\x0c', 'X.-M. Hou, K.-C. Chou / Journal of Alloys and Compounds 509 (2011) 2395-2400  2397  Fig. 1. The isothermal oxidation curves of ZrC powder at different oxygen pressures (a) 1.3 × 10−3 MPa, (b) 2.6 × 10−3 MPa, and (c) 7.9 × 10−3 MPa.  Fig. 2. Oxidation of ZrC powder at different oxygen pressures (a) non-isothermal oxidation and (b) isothermal oxidation.  \\x0c', '2398  X.-M. Hou, K.-C. Chou / Journal of Alloys and Compounds 509 (2011) 2395-2400  0  0  O2  O2  \\x01Ed represents the apparent activation energy of diffusion; BdT is a function of PO2 , P eq and R0 , in which P eq is the oxygen partial pressure in equilibrium with oxide and should be related to temperature T. K 0ˇ and D0ˇ are constants independent of temperature but replying on the property of the material; vm is a coefficient related to the density of reactant and product; R0 is the radius of the particle. If the value of P eq is very small or the temperature coefficient of P eq  can be neglected, thus BdT will be constant as the oxygen partial pressure and the particle radius are fixed. According to Eq. (2), two parameters, i.e., \\x01Ed and BdT should be extracted to describe the oxidation behavior, which could be fitted from the experimental data by regression method. Therefore, the oxidation kinetics formula can be described as follows: = 1.3 × 10 −3 MPa (675-737 K)  O2  O2  PO2  (cid:7)  (cid:8)  (cid:2) = 1 −  1 − 86.07 exp  − 7679.82  T  (cid:9) √  (cid:10)3  t  (4)  Fig. 3. The isothermal oxidation of ZrC pellet.  3.1.2. Oxidation kinetics of ZrC pellet  Since the oxidation of ZrC pellet investigated by Kuriakose and Margrave was controlled by chemical reaction [11], the linear rate law was employed and the result can be well fitted with the activation energy to be 72 kJ/mol.  3.2. Application of Chou model  3.2.1. Oxidation kinetics of ZrC powder  In view of the isothermal oxidation behavior of ZrC powder at the oxygen pressure of 1.3 × 10−3 MPa, its kinetics was controlled by diffusion [7]. One formulae of Chou’s model can be employed to treat it, i.e.,  (cid:2)  (cid:3)  (cid:2) = 1 −  1 −  where  BdT =  (2K 0ˇ  0  D0ˇ 0  (cid:4)3  exp(−\\x01Ed /RT )t  BdT  (cid:5)  1  /vm )((  −  PO2  (cid:6)  P eq O2  )/R2 0 )  (2)  (3)  Similarly, the isothermal oxidation behavior of ZrC powder at the oxygen pressure of 2.6 × 10−3 and 7.9 × 10−3 MPa, respectively can also be deduced. Because the oxidation mechanism changes at higher temperature, i.e., 743 K, thus the oxidation behavior should be treated, respectively. The equations are following: = 2.6 × 10 −3 MPa  PO2  (cid:8)  (cid:8)  (cid:8)  (cid:8)  (cid:10)3  (cid:10)3  (cid:9) √  t  (cid:9) √  t  (cid:10)3  (cid:10)3  (cid:9) √  t  (cid:9) √  t  (5)  (6)  (7)  (8)  (cid:7)  (cid:7)  (cid:7)  (cid:7)  (a) 663-743 K  (b) 758-803 K  (cid:2) = 1 −  1 − 102.6 exp  − 7679.82  T  (cid:2) = 1 −  1 − 2582.0 exp  − 9953.1  T  = 7.9 × 10 −3 MPa  PO2  (a) 652-743 K  (cid:2) = 1 −  1 − 106.6 exp  − 7679.82  T  (b) 753-800 K  (cid:2) = 1 −  1 − 3162.3 exp  − 9953.1  T  Fig. 4. A comparison of plots from Chou model and experimental data for isothermal oxidation of ZrC powder at an oxygen pressure of 2.6 × 10−3 MPa (a) 663-730 K and (b) 758-803 K.  \\x0c', 'X.-M. Hou, K.-C. Chou / Journal of Alloys and Compounds 509 (2011) 2395-2400  2399  Fig. 5. A comparison of plots from Chou model and experimental data for isothermal oxidation of ZrC powder at an oxygen pressure of 7.9 × 10−3 MPa (a) 652-739 K and (b) 753-800 K.  The curves obtained from the above equations are also shown in the figures for comparison (as shown in Figs. 1a, 4a and b and 5a and b, respectively) and got a good agreement. Shimada and Ishii [7] also investigated the oxygen pressure on oxidation of ZrC powder under both non-isothermal and isothermal conditions. Here the following equations are employed to treat the oxidation behavior, i.e., Isothermal oxidation:  (cid:2) = 1 −  ⎛ ⎜⎜⎝1 −  (cid:14)(cid:15)(cid:15)(cid:16) ( (cid:5)  PO2  −  (cid:6)  P eq O2  )t  Bp  ⎞ ⎟⎟⎠  3  (9)  where  Bp =  1 /vm )(exp(−\\x01Ed /RT )/R2 0 )  (2K 0ˇ  0  D0ˇ 0  (10)  Non-isothermal oxidation:  (cid:2) = 1 −  ⎛ ⎜⎜⎝1 −  (cid:14)(cid:15)(cid:15)(cid:16) ( (cid:5)  PO2  −  (cid:6)  P eq O2  )exp(−\\x01Ed /RT )  Bcp  T − T0 \\x01  ⎞ ⎟⎟⎠  3  (11)  where Bcp = 1/(2K 0ˇ rate. According to Eq. (11), the oxidation fraction of ZrC powder as a function of oxygen pressure and heating rate under non-isothermal condition is as follows:  0  /vm D0ˇ  0  /R2  0 ) and \\x01 is temperature-increasing  (cid:2) = 1−  (cid:20)  1−5345.2 exp  (cid:8)  − 9766.7  T  (cid:9) (cid:6)  (T −298)(  (cid:5)  PO2  −0.0241)  (cid:21)3  (12)  To compare with the experimental data, the results calculated at the oxygen pressure of 1.58 × 10−2 and 3.95 × 10−2 MPa are listed in Fig. 2a. The coincidence of the theoretical calculation results with the experimental data indicates that this theory is reasonable. The effect of oxygen pressure on oxidation under isothermal condition can also be quantitatively discussed using Eq. (9). The experimental oxidation data at 753 K with the oxygen pressure of 1.3 × 10−3 and 7.9 × 10−3 MPa were used to testify Eq. (9) and the  result showed that a good agreement was obtained (Fig. 2b). The formula is as follows:  (cid:2) = 1 −  (cid:20)  1 − 0.019  (cid:6)  (  (cid:5)  PO2  − 0.00132)t  (cid:21)3  (13)  3.2.2. Oxidation kinetics of ZrC pellet  In view of the oxidation of ZrC pellet, its oxidation was controlled by chemical reaction [11]. Chou’s model has also developed a series of formulae to deal with this condition as following:  (cid:2) =  1  (cid:6)T  exp  (cid:8)  − \\x01Er RT  (cid:9)  t  (14)  where  (cid:6)T =  vm L0  k0 (  (cid:5)  pO2  −  (cid:6)  peq O2  )  (15)  and \\x01Er represents the apparent activation energy of chemical reaction, P eq is oxygen partial pressure equilibrium with oxide, k0 is a temperature-independent constant; vm is a coefficient related to the density of reactant and product, L0 is the original thickness of the pellet. If all these parameters are known, of course, one can find the relation between reacted fractions (cid:2) with time t. In some particular cases, one can combine part of these parameters to construct an auxiliary function like (cid:6)T (see Eq. (15)) that can be found through curve fitting. Eq. (14) is employed to deal with the oxidation behavior of ZrC pellet and the isothermal oxidation equation is given as follows:  O2  \\x01m  A  = 0.379 exp  (cid:8)  − 9965.12  T  (cid:9)  t  (16)  Based on the above formulae, a set of theoretical oxidation curves for ZrC pellet is also plotted in Fig. 3 for comparison. A fair agreement has been gotten and the corresponding activation energy is 82.9 kJ/mol.  3.2.3.  Comparison with previous oxidation studies of ZrC  The experimental data offered by Shimada and Ishii [7] and Kuriakose et al. [11] showed that the oxidation mechanism of ZrC powder seems not a simple one. The effects of temperature and oxygen pressure on the oxidation behavior of ZrC were evident. As mentioned above, the oxidation reaction of ZrC is a kind of complicated heterogenous reaction. Most researchers studied the  \\x0c', '2400  X.-M. Hou, K.-C. Chou / Journal of Alloys and Compounds 509 (2011) 2395-2400  Table 1 Comparison of the calculation error from the model used in the literature and the new model.  Data  \\x01 (average relative error)  Calculated from the models used in the literature (%)  Calculated from Chou model (%)  Oxidation of ZrC offered by Shimada and Ishii [7] Oxidized with the oxygen pressure of 1.3 × 10−3 MPa Oxidized with the oxygen pressure of 2.6 × 10−3 MPa 663-743 K 758-803 K Oxidized with the oxygen pressure of 7.9 × 10−3 MPa 652-743 K 753-800 K Oxidation of ZrC pellet offered by Kuriakose and Margrave [11]  13.9  15.8 22.9  12.5 24.0 9.5  6.7  8.1 7.6  7.3 10.7 7.2  oxidation kinetics of ZrC materials using the linear rate law or Jander’s equation. The major difference between Chou’s model and the models used in the literatures is that, firstly, Chou’s model can express the relation between reacted fraction and temperature, oxygen partial pressure and many other variables explicitly. Secondly, Chou’s model has revealed the physical meaning of the parameter “k” of the models used in the literature, which can be expressed as a function of temperature T, particle size R0 , oxygen pressure PO2 , etc. All these parameters appearing in Chou’s model have clear physical meaning because the model has been derived based on a real physical picture without vagueness assumption. Therefore, it is impossible to assign an arbitrary value to them if that does not meet the requirement of physical meaning. Then from the view of mathematical treatment, Chou’s model is much better than the models used in the literatures in extracting the activation energy because it only needs to perform regression just once. Therefore, the calculated error from Chou’s model should be smaller. The calculated error from the new model and the model used in the literatures can be obtained by the following equation:  \\x01 = 1  N  |((cid:5)i )mea − ((cid:5)i )cal | |((cid:5)i )mea |  × 100%  (17)  N(cid:22)  i=1  ((cid:5)i )cal  where \\x01 is the average relative error, ((cid:5)i )mea is experimental data, is the value calculated from the model and N is the number of experimental points. Table 1 shows the calculation results from the model used in the literatures and the new model presented here in treating the oxidation of ZrC materials from which it can be seen that the error calculated by Chou’s model is smaller than that calculated by the model used in the literatures. Therefore it makes us feel more comfortable to add some predicted lines within the same reaction mechanism that might be interested in application.  4. Conclusion  The effects of temperature and oxygen pressure on oxidation behavior of ZrC materials have been studied from theoretical aspect based on the experimental data. The experimental results showed that the effect of temperature and oxygen pressure on the oxidation behavior of ZrC was evident. Two sets of experimental data were used in this analysis: Shimada et al. and Kuriakose et al. For Shimada’s data, although the oxidation kinetics was diffusion controlled during the whole exper imental temperature range, the activation energy was different because of the phase transition. The oxidation behavior should be described by different equations, i.e., Eqs. (4)-(8). In view of the effect of oxygen pressure on oxidation behavior under non-isothermal and isothermal conditions, we also deduced two quantitative formulae, Eqs. (12) and (13) and got a good agreement with experimental data. Concerning the oxidation of ZrC pellet investigated by Kuriakose et al., the oxidation behavior followed a linear rate law and the oxidation kinetics can be expressed as follows:  (cid:8)  (cid:9)  \\x01m  A  = 0.379 exp  − 9965.12  T  t  Based on the error analysis, it is clear that Chou’s model can give more accurate results than that offered by the traditional models used in the literature.  Acknowledgements  The authors sincerely thank the support from Key Lab. of Ecologic & Recycle Metallurgy, Ministry of Education, University of Science and Technology Beijing. The authors also would like to express their thanks to the National Natural Science Foundation of China (Nos. 50974084 and 50874013).  References  [1] A. Bellucci, D. Gozzi, T. Kimura, T. Noda, S. Otani, Surf. Coat. Technol. 197 (2005) 294-302. [2] D. Ferro, S.M. Barinov, J.V. Rau, A. Latini, R. Scandurra, B. Brunetti, Surf. Coat. Technol. 200 (2006) 4701-4707. [3] Y. Ozaki, R.H. Zee, Mater. Sci. Eng. A 202 (1995) 134-141. [4] C. Kral, W. Lengauer, D. Rafaja, P. Ettmayer, J. Alloys Compd. 265 (1998) 215-233. [5] S. Shimada, M. Nishisako, M. Inagaki, K. Yamamoto, J. Am. Ceram. Soc. 78 (1995) 41-48. [6] R.W. Bartlett, M.E. Wadsworth, I.B. Cutler, AIME Trans. Metall. Soc. 227 (1963) 467-472. [7] S. Shimada, T. Ishii, J. Am. Ceram. Soc. 73 (1990) 2804-2808. [8] G.A. Rama Rao, V. Venugopal, J. Alloys Compd. 206 (1994) 237-242. [9] S. Shimada, Solid State Ionics 149 (2002) 319-326. [10] S. Shimada, M. Inagaki, M. Suzuki, J. Mater. Res. 11 (1996) 2594-2597. [11] A.K. Kuriakose, J.L. Margrave, J. Electrochem. Soc. 111 (1964) 827-831. [12] S. Shimada, Solid State Ionics 101-103 (1997) 749-753. [13] S. Shimada, Solid State Ionics 141-142 (2001) 99-104. [14] K.C. Chou, X.M. Hou, J. Am. Ceram. Soc. 92 (2009) 585-594. [15] K.C. Chou, J. Am. Ceram. Soc. 89 (2006) 1568-1576. [16] X.M. Hou, K.C. Chou, X.J. Hu, H.L. Zhao, J. Alloys Compd. 459 (2008) 123-129.  \\x0c']"
},{
  "_id": 114,
  "PDF": "Investigations on synthesis of ZrB2 and development of new composites with HfB2 and TiSi2.pdf",
  "Text": "['Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Contents lists available at ScienceDirect  Int.  Journal of Refractory Metals and Hard Materials  j o u r n a l h o m e p a g e : w w w. e l s e v i e r . c o m / l o c a t e / I J R M H M  Investigations on synthesis of ZrB2 and development of new composites with HfB2 and TiSi2  J.K. Sonber ⁎, T.S.R. Ch. Murthy, C. Subramanian, Sunil Kumar 1, R.K. Fotedar, A.K. Suri  Materials group, Bhabha Atomic Research Centre, Mumbai,  India-400 085  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 26 February 2010 Accepted 10 June 2010  Keywords:  ZrB2 TiSi2 Composite Synthesis Densification Microstructure Oxidation studies  This paper presents the results of experimental investigations carried out on the synthesis of pure ZrB2 by boron carbide reduction of ZrO2 and densification with the addition of HfB2 and TiSi2. Process parameters and charge composition were optimized to obtain pure ZrB2 powder. Monolithic ZrB2 was hot pressed to full density and characterized. Effects of HfB2 and TiSi2 addition on densification and properties of ZrB2 composites were studied. Four compositions namely monolithic ZrB2, ZrB2 + 10% TiSi2, ZrB2 + 10% TiSi2 + 10% HfB2 and ZrB2 + 10% TiSi2 + 20% HfB2 were prepared by hot pressing. Near theoretical density (99.8%) was obtained in the case of monolithic ZrB2 by hot pressing at 1850 °C and 35 MPa. Addition of 10 wt.% TiSi2 resulted in an equally high density of 98.9% at a lower temperature (1650 °C) and pressure (20 MPa). Similar densities were obtained for ZrB2 + HfB2 mixtures also with TiSi2 under similar conditions. The hardness of monolithic ZrB2 was measured as 23.95 GPa which decreased to 19.45 GPa on addition of 10% TiSi2. With the addition of 10% HfB2 to this composition, the hardness increased to 23.08 GPa, close to that of monolithic ZrB2. Increase of HfB2 content to 20% did not change the hardness value. Fracture toughness of monolithic sample was measured as 3.31 MPa m1/2, which increased to 6.36 MPa m1/2 on addition of 10% TiSi2. With 10% HfB2 addition the value of KIC was measured as 6.44 MPa m1/2, which further improved to 6.59 MPa m1/2 with higher addition of HfB2 (20%). Fracture surface of the dense bodies was examined by scanning electron microscope. Intergranular fracture was found to be a predominant mode in all the samples. Crack propagation in composites has shown considerable deflection indicating high fracture toughness. An oxidation study of ZrB2 composites was carried out at 900 °C in air for 64 h. Specific weight gain vs time plot was obtained and the oxidized surface was examined by XRD and SEM. ZrB2 composites have shown a much better resistance to oxidation as compared to monolithic ZrB2. A protective glassy layer was seen on the oxidized surfaces of the composites.  © 2010 Elsevier Ltd. All rights reserved.  1. Introduction  Zirconium diboride is a leading material in the category of ultra high temperature ceramics (UHTC) due to very high melting point (3245 °C), (57.9 WM−1 K−1), high thermal conductivity good thermal shock resistance, low coefficient of thermal expansion (5.9 × 10−6 °C−1), retention of strength at elevated temperatures and stability in extreme environments [1-3]. ZrB2 is considered a candidate material for hypersonic flight, atmospheric re-entry and rocket propulsion [1,4,5]. It gets wetted but not attacked by molten metals and hence is used for holding molten metal and as thermo-well tubes in metal processing [6]. Good electrical conductivity makes it suitable for electrode application in Hall-Heroult cell and electric discharge machining [7-9]. ZrB2 can be synthesized by (a) reaction between Zr and B [10] (b) borothermic reduction of ZrO2 [11], (c) boron carbide reduction of ZrO2 in the presence of carbon [12], (d) carbothermic reduction of ZrO2 and  ⁎ Corresponding author. Tel.: +91 22 2559 0473; fax: +91 22 2550 5151. E-mail address: jitendra@barc.gov.in (J.K. Sonber). 1 Post Irradiation Examination Division, BARC.  0263-4368/$ - see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2010.06.007  B2O3 [1] (e) metallothermic reduction of ZrO2 and B2O3 [13-15] and (f) chemical vapor deposition. The synthesis of ZrB2 from its elements is uneconomical due to the high cost of both Zr and B powders. Borothermic reduction of ZrO2 to obtain ZrB2 also involves the use of expensive boron powder. In carbothermic reduction of ZrO2 and B2O3, loss of boron occurs due to the evaporation of boron oxides during the reaction and the product is not pure ZrB2 but a mixture of borides. In case of metallothermic reduction of ZrO2 and B2O3 the product gets contaminated with the metal borides of reductant metal. ZrB2 powder can also be synthesized by SHS reaction between ZrO2, B2O3 and Mg. Chemical vapor deposition techniques are suitable for coating and not for bulk production of powders. Preparation of ZrB2 by boron carbide reduction of ZrO2 [12] according to reaction (1) seems to be the best route since it involve use of cheap raw materials and results in pure boride with minimum boron loss.  ZrO2 þ 1=2B4C þ 3=2C→ZrB2 þ 2CO↑  ð1Þ  Despite having excellent properties, the actual uses of ZrB2 are limited due to its poor sinterability and low fracture toughness. Due to  \\x0c', \"22  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  strong covalent bonding and low self-diffusion, high temperature and external pressure are required to densify monolithic ZrB2 [1,10]. Martinez et al. [16] have densified monolithic ZrB2 to 86.5% ρth by hot pressing at 1900 °C and 30 MPa pressure. Chamberlain et al. [17] have obtained a density of  98% ρth by pressureless sintering ZrB2 at 2150 °C for 540 min. The same authors have reported the densification of ZrB2 powder to 99.8% by hot pressing at 1900 °C and 32 MPa [18]. By reactive hot pressing of zirconium and boron powder mixture, 99% dense ZrB2 was obtained at 2100 °C [19]. Several additives have been tried to improve the sinterability and properties of ZrB2. Mishra et al. [20] have reported that oxygen present on the surface of ZrB2 hinders densification. By the addition of TiC and C they have achieved 94% ρth by pressureless sintering at 1800 °C in Ar atmosphere. Zhu et al. [21] have sintered ZrB2 powder to near theoretical density at 1900 °C without the application of external pressure using carbon coated starting powders. Carbon addition was found to effectively remove oxide impurities from the surface of ZrB2 particles. Fahrenholtz et al. [22] have sintered ZrB2 to full density at 1850 °C using a combination of B4C and carbon as additive. Rangaraj et al. [23] have obtained ZrB2-ZrC composite by reactive hot pressing of Zr and boron carbide mixture at 1600 °C and 40 MPa. Apart from these additives, silicon carbide is the most common, as it has been found to enhance the strength and oxidation resistance of zirconium diboride ceramics [1,18]. Yan et al. [24] have densified ZrB2-20 wt.% SiC by pressureless sintering at 2250 °C using 4 wt.% Mo as additive. The composite showed high fracture toughness of 5.39 MPa m1/2 .. Zhu et al. [25] have hot pressed ZrB2-30 vol.% SiC composite to 99.8% ρth at 1900 °C and 32 MPa. The composite exhibits high flexural strength of 900 MPa. M. Zhu et al. [26] have used liquid polycarbosilane (LPCS) as additive to obtain 95% dense ZrB2-SiC composite by pressureless sintering at 1900 °C. Addition of 4% Ni improves the densification and results in 98% density while hot pressing at 1850 °C [16,27] . Monteverde et al. [28] have found that addition of 2.5% Si3N4 causes formation of grain boundary glassy phase and results in 98% ρth density at 1700 °C and 30 MPa. MoSi2 (15 to 20%) was found to help in densification, resulting in 98% dense ZrB2 at 1750 °C and 30 MPa [29,30]. Sciti et al. [31] have densified ZrB2-20 vol.% MoSi2 at 1850 °C and obtained fine grained structure. Recently Shu-Qi Guo et al. [32] have studied the effect of ZrSi2 addition on pressureless sintering of ZrB2 and obtained 99.5% ρth sample at 1600 °C in vacuum by the addition of 20% ZrSi2. As per authors' knowledge there is no literature on the use of TiSi2 and HfB2 as sinter additive to ZrB2. This paper presents the study on  Fig. 2. Thermogravimetry of ZrB2 synthesis.  the effect of TiSi2 and HfB2 addition on processing and properties of ZrB2.  2. Experimental  2.1. Starting material  Raw materials used were ZrO2 (99% purity; 8.34 μm median diameter), boron carbide powder (78.5% B, 19.5% C, b 1% O, 0.02% Fe, 0.02% Si, 5.34 μm median diameter; supplied by M/S Boron 13.9 μm median Carbide India) and petroleum coke (C-99.4%, diameter, supplied by M/S Assam carbon, India). All the raw  Fig. 3. XRD Pattern of thermogravimetry product.  Table 1 Effect of temperature and charge composition.  S. no.  Molar ratio ZrO2:B4C:C  Temperature (°C)  Weight loss (%)  Phases present  Carbon (%)  Oxygen (%)  1. 2. 3. 4. 5. 6. 7.  2:1:3 2:1:3 2:1:3 2:1:3 2:1:3 2:1.1:2.7 2:1.1:2.7  1200 1400 1600 1700 1800 1800 1875  18.47 27.15 30.28 31.66 33.04 32.67 33.01  ZrB2, ZrO2 ZrB2, ZrO2 ZrB2, ZrB, ZrO2 ,C ZrB2, ZrB, C ZrB2, ZrB, C ZrB2 ZrB2  9.7 8.2 7.3 6.1 3.2 1.3 0.06  5.22 2.9 2.4 2 0.57 1.5 0.5  Fig. 1. XRD pattern of starting materials ZrO2, B4C and petroleum coke.  \\x0c\", \"J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  23  2.2. Synthesis  A small pellet of the reactants in stoichiometric quantity (as per reaction 11) was prepared by weighing accurate quantity of individual components, mixing it thoroughly and pelletizing. This pellet was used for thermogravimetric investigations (Setaram TAG 24) up to 1500 °C. For regular synthesis of ZrB2, weighed quantities of zirconium dioxide, boron carbide and petroleum coke in various ratios were mixed thoroughly in motorized mortar and pestle. The powder mixture was then pelletized under pressure of about 280 MPa to obtain pellets of  Fig. 4. Free energy change of reactions (thermodynamic data from Barin's table [34]).  materials were dried in an oven at 100 °C to remove moisture content be fore use . XRD pattern of the raw materials are presented in Fig. 1.  Fig. 5. XRD pattern of the products obtained by varying charge composition. MR: molar ratio (ZrO2:B4C:C).  Fig. 6. Effect of temperature and charge composition on carbon and oxygen content of the product.  Fig. 7. SEM image of powders (a) ZrB2, (b) HfB2 and (c) TiSi2.  \\x0c\", \"24  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Fig. 8. Particle size distribution of powders (a) ZrB2, (b) HfB2, (c) TiSi2 and (d) mixed charge.  20 mm diameter. While mixing and pelletizing, all the materials coming in contact with the powder were made of tungsten carbide. The pellets were then charged in a graphite crucible and heated in an induction furnace under a dynamic vacuum of 2 × 10−5 mbar and at fixed temperature between 1200 and 1875 °C for 2 h. Temperature of the charge was measured using a two-colour pyrometer with an accuracy of ± 22 °C. After completion of the reaction, the furnace was cooled to room temperature in vacuum and the reacted pellets were taken out, crushed and ground to fine size using high-energy cup grinding mill with tungsten carbide lining. The major phases of the powders were identified by XRD (Cu Kα (λ = 1.5404 Å radiation in a Philips PW1830 diffractometer) and impurities were analyzed by chemical methods. The median particle diameter and particle size distribution were measured using laser particle size analyzer (CILAS PSA 1064 L). SEM (20 kV, Philips, FEI XL30) of the powders was also used to cross check the size and morphology.  2.3. Densification and characterization  For densification, weighed quantities of fine zirconium diboride, hafnium diboride and titanium disilicide were mixed thoroughly using a motorized mortar and pestle in dry condition for 1 h and powder mixture of four different compositions were prepared as (1) ZrB2 (2) ZrB2 + 10% TiSi2, (3) ZrB2 + 10% TiSi2 + 10% HfB2 (4) ZrB2 + 10% TiSi2 +  Table 2 Effect of sinter additives on densification and properties of ZrB2.  Additive  Temp. (°C)  Pressure (MPa)  Density (%)  Hardness (GPa)  KIC  (MPa m1/2)  nil 10 wt.% TiSi2 10% HfB2 + 10 wt.% TiSi2 20% HfB2 + 10 wt.% TiSi2  1850 1650 1650 1650  35 20 20 20  99.8 98.9 99.6 98.4  23.91 ± 1.5 19.45 ± 1.9 23.08 ± 1.3 23.66 ± 1.8  3.31 ± 0.2 6.36 ± 1.0 6.44 ± 1.1 6.59 ± 0.7  20% HfB2. The powders were then loaded in a high density graphite die (12 mm hole) and hot pressed at temperatures of 1650 °C to 1850 °C under a pressure of 20 to 35 MPa for 60 min in a high vacuum (1 × 10−5 mbar) chamber. The pellets were ejected from the die after cooling and the density measured by Archimedes' principle. Densified samples were polished to mirror finish using diamond powder of various grades from 15 to 0.25 μm in an auto polisher (laboforce-3, Struers). Microhardness was measured on the polished surface at a load of 100 g and dwell time of 10 s. The indentation fracture toughness (KIC) data were evaluated by crack length measurement of the crack pattern formed around Vickers indents (using 10 kg load), adopting the model formulation proposed by Anstis et al. [33] KIC = 0.016(E/H)1/2P/c3/2, where E is the Young's modulus, H the Vickers hardness, P the applied indentation load, and c  Fig. 9. XRD Pattern of ZrB2 based composites.  \\x0c\", 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  25  Table 3 Oxidation parameters for ZrB2 composites.  Sample  Monolithic ZrB2 ZrB2 + 10 wt.% TiSi2 ZrB2 + 10% HfB2 + 10 wt.% TiSi2 ZrB2 + 20% HfB2 + 10 wt.% TiSi2  m  0.97 2.52 2.38 2.02  Km  4.5 × 10−7 1.96 × 10−10 1.67 × 10−10 2.03 × 10−10  Kp (Kg m−4 s−1) × 108  1.25 × 10−8 1.80 × 10−9 9.53 × 10−10 2.22 × 10−10  the half crack length. The reported values of hardness and fracture toughness are the average of five measured values. Fractured surface of dense pel lets was analyzed by scanning electron microscope.  2.4. Oxidation  Hot pressed pellet of diameter 12 mm was cut into thin slice of 3 mm thickness by high speed diamond cutter. All the surfaces of the cut sample were polished with emery papers (1/0, 2/0, 3/0, 4/0) and finally with diamond paste up to 1 μm finish. Oxidation tests were conducted in a resistance heated furnace. In order to avoid oxidation during heating, the sample was directly inserted into the furnace after the furnace temperature reached 900 °C. Samples were placed in an alumina crucible kept into the furnace. The sample was oxidized for different time intervals (0.5, 1, 2, 4, 8, 16, 32, and 64 h) at 900 °C. The sample was carefully weighed before and after exposure, to determine the weight change during the oxidation process. The oxidation products were identified using XRD. The morphology and nature of oxide layer were understood by observing the surface in a scanning electron microscope (SEM).  3. Results and discussion  3.1. Synthesis  Fig. 10. SEM image of ZrB2 + 10%TiSi2 + 20%HfB2 composite and elemental analysis of different phases.  Fig. 2 presents the weight loss vs. time curve obtained by thermogravimetric experiment of stoichiometric charge mixture as per reaction 1. The reaction starts at 1200 °C and results in a total  Fig. 11. Elemental mapping of Zr, Ti and Si  in ZrB2 + 10%TiSi2 + 20%HfB2 composite.  \\x0c', '26  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  loss of  27.5%. The weight weight loss observed is lower than the theoretical weight loss of 33.0%. The XRD pattern revealed the presence of ZrB2, ZrB and graphite phases (Fig. 3). In the case of synthesis experiments carried out in an induction furnace at 1200 °C, the weight loss was 18.47% and the product was impure, containing ZrB2 and ZrO2 phases (Table 1). The carbon and oxygen contents were 9.7 and 5.22% respectively. At 1400 °C, the weight loss increased to 27.15% and carbon and oxygen contents were less by 8.2 and 2.9%. At 1600 °C, weight loss increases to 30.28% but the product contained ZrB2, ZrB, ZrO2 and C. At temperatures of 1700 °C and above, ZrO2 was absent in the product but ZrB was seen. At 1800 °C the weight loss was close to a theoretical value of 33.0%, however the product was  Fig. 12. Fracture surfaces of (a) monolithic ZrB2 (b) ZrB2 + 10%TiSi2 (c) ZrB2 + 10%TiSi2 + 10%HfB2.  composed of ZrB2, ZrB, and carbon. The presence of boron deficient phase (ZrB) even after the treatment at 1800 °C indicates the loss of boron from the charge, which could occur by following reaction.  ZrO2 þ 3=4B4C ¼ ZrB2 þ 1=2B2O3 þ 1=2CO þ 1=4C  ð2Þ  The thermodynamic calculation [34] (Fig. 4) indicates that both reactions (1) and (2) could take place on heating and hence the loss of boron by evaporation as B2O3 during the reaction. Hence to obtain a single-phase ZrB2 phase, it is necessary to add an excess boron in the charge. Fig. 5 presents the XRD patterns of product obtained by stoichiometric and modified charge mixture. It evidently tells that ZrB and graphite phase are not present and single phase ZrB2 is obtained by reacting the modified charge mixture (molar ratio, ZrO2:B4C:C, 2:1.1:2.7) at 1800 °C. The oxygen and carbon contents were found to be 1.5 and 1.3% respectively. This product was further purified to low levels of oxygen (0.06%) and carbon (0.5%) by heating to 1875 °C and soaking at this temperature in vacuum for 30 min. The progressive removal of carbon and oxygen from the product with reaction temperature and the final purification using modified charge are presented in Fig. 6. Hong Zhao et al. [12] have also reported the intermediate reaction resulting in the loss of boron as B2O3.  3.2. Densification and characterization  SEM images (Fig. 7) of the powder shows that ZrB2 and HfB2 particles are of  2-3 μm whereas TiSi2 particles are large at 15-20 μm size. Fig. 8 presents the particle size distribution of ZrB2, HfB2, TiSi2 and mixed charge. It shows that particle size distribution is monomodal for ZrB2 and HfB2 whereas bimodal for TiSi2. During mixing, size reduction of the particles also takes place and the particle size distribution of the mixed powders is narrow and monomodal. Median diameter of the mixed charge was measured as 2.67 μm. Table 2 presents the effects of hot pressing parameters on the density and mechanical properties of the pellets. In the case of monolithic ZrB2, a near full density (99.8% TD) was obtained at 1850 °C with a pressure of 35 MPa and dwell time of 60 min. Addition of 10 wt.% TiSi2 resulted in densification of 98.9% TD at a relatively low tempera ture o f 1650 °C and a low pressure o f 20 MPa . Composites with 10 and 20% HfB2 content were also hot pressed to nearly full density at 1650 °C and 20 MPa . The enhanced sintering at lower hot pressing temperature and pressure is due to the liquid phase sintering with low melting (1540 °C) additive TiSi2 and the low melting reaction product ZrSi2 (1620 °C). XRD pattern of the dense pellets are shown in Fig. 9. All the three samples indicate the presence of crystalline ZrB2 and ZrSi2. ZrSi2 is formed during sintering by the following reaction.  ZrB2 þ TiSi2→ZrSi2 þ TiB2  ð3Þ  TiB2 is not seen in the XRD pattern of the sintered product, probably due to the formation of ZrB2-TiB2 solid solution. HfB2 also forms a solid solution with ZrB2 and hence not seen as a distinct phase in XRD pattern. Post et al. [35] have reported that, these borides have complete mutual solubility. Compar ison o f o ther s in ter add i t ives f rom l i tera ture is discussed below. The addition of 20% MoSi2 results in 98.1% TD on hot pressing at 1800 °C and 30 MPa [30]. 20% SiC addition resulted in near theoretical density on hot pressing at 2000 °C and 30 MPa [36]. Guo et al. [37] have reported that addition of 5% Re2O3 (Re = Y,Yb,La,Nd) to ZrB2-20% SiC results in N 99% TD on hot pressing at 1900 °C. Wang et al. [38] have reported that 10% Mo addition gives a density of 98.9% TD on hot pressing at 1950 °C and 20 MPa . Zhu et al. [39] have reported that addition of 3-10%  \\x0c', 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  27  Fig. 13. Crack propagation in (a) monolithic ZrB2 (b) ZrB2 + 10%TiSi2 (c) ZrB2 + 10%TiSi2 + 10%HfB2 (d) ZrB2 + 10%TiSi2 + 20%HfB2.  Al2O3 and Y2O3 and 20% SiCw to ZrB2 gives a density N 97% TD on hot pressing at 1800 °C. From the above it is clear that the addition of TiSi2 in the present investigation is found very effective in lowering the sintering temperature to 1650 °C. The addition of TiSi2 to TiB2 and HfB2 has also shown similar results [40,41].  3.3. Mechanical properties  Variations in Vickers hardness and fracture toughness of ZrB2 composites are presented in Table 3. Hardness of monolithic sample was measured as 23.91 GPa, which reduced to 19.45 GPa with the addition of 10 wt.% TiSi2. The decrease in hardness is due to the lower hardness of TiSi2 (8-10 GPa) [40]. The effect of TiSi2 is nullified with HfB2 addition. The hardness value composite with 10% HfB2 addition was 23.08 GPa, which marginally increased further to 23.66 GPa with 20 wt.% HfB2. The increase in hardness could be attributed to (Zr-Hf)B2 solid solution formation. Hardness of monolithic ZrB2 has been reported to be in the range of 22-23 GPa [1,2,18,42]. Akin et al. [43] have reported the hardness of ZrB2-20-60% SiC composite to be 26 GPa. Fracture toughness of monolithic ZrB2 sample was measured as 3.31 MPa m1/2, which increased considerably to 6.36 MPa m1/2 with 10% TiSi2 addition. With further addition of HfB2, the fracture toughness of ZrB2 + T iS i2 + H fB2 compos ite was found to be marginally high at 6.44-6.59 MPa m1/2. Fracture toughness of monolithic ZrB2 has been reported as 3.5 MPa m1/2[18]. Enhancement of fracture toughness from 4.52 to 7.98 MPa m1/2 by 5% Mo addition is reported by Wang et al. [38].Zhu et al. [39] have reported the fracture toughness of 6.7 MPa m1/2 for ZrB2 + 20% SiCw + 3%  YAG. Similar value of fracture toughness was reported for ZrB2-SiC nanocomposite prepared by using SiC nanopowder [44].  3.4. Microstructure evolution  Fig. 10 presents the secondary electron image of ZrB2 + 10% TiSi2 + 20% HfB2 composites. It shows the presence of dark phase in the gray matrix. EDS pattern of the phases are also inserted into the picture.  Fig. 14. Specific weight gain vs time plot for ZrB2 based composites. (Temp: 900 °C).  \\x0c', '28  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Gray matrix is essentially ZrB2 in which Ti and Hf have diffused whereas dark phase contains mainly Si with some Zr. Fig. 11 presents the elemental mapping for Zr, Ti and Si. It shows that Zr and Ti are distributed uniformly whereas Si is present only in the black phase. Fig. 12 (a-c) presents the fracture surfaces of monolithic ZrB2, ZrB2 + 10% TiSi2 and ZrB2 + 10% TiSi2 + 10% HfB2 composites. The mode of fracture is seen to be intergranular in all the samples. Regular faceted grains are visible. Fig. 13 (a) presents the features of indentation crack in monolithic ZrB2 and Fig. 13 (b-d) in composites. The crack propagation line is almost straight in monolithic ZrB2 whereas considerable deflections are observed in the composite samples, which explain the high fracture toughness.  3.5. Oxidation study  The weight gain data obtained during oxidation at 900 °C as a function of time for ZrB2 composites are presented in Fig. 14. Continuous weight gain with time is observed in all the samples. In the case of composites, the rate of oxidation was found to decrease with increase in time which indicates the formation of protective layer. In monolithic ZrB2, oxidation rate was found to be constant. In order to understand the nature of oxidation, the oxidation data was fitted in the general rate equation for all the composites.  ð  Þm = Km :t Δw = A  ð4Þ  where Δw is the change in weight, A—surface area of the sample, t— oxidation time and Km—rate constant. Km and m values are presented in Table 3. The composite samples have shown a parabolic pattern of oxidation whereas monolithic has shown linear oxidation. Better oxidation resistance of the composite samples is attributed to the formation of silica based glassy layer. The SEM microstructures (Fig. 15) of oxidized surfaces evidently show the formation of protective glassy phase. The glassy phase was analyzed by EDS to contain mainly silicon ( 46 at.%) and oxygen ( 52 at.%). Zirconium ( 0.8 at.%) and titanium ( 0.2 at.%) are also present in very small quantity. A typical EDS pattern from the oxidized surface is shown in Fig. 16. Formation of protective borosilicate glass has been reported in literature [45,46]. The glass formed in this study could be borosilicate glass. EDS analysis has some limitations for detection of boron. Fig. 17 presents the XRD pattern of the oxidized surface. The major crystalline phase in all the composites is confirmed as ZrO2. Peaks of ZrSiO4 and TiO2 are also present. The following reactions are possible during the oxidation process.  2=5ZrB2 þ O2→2=5ZrO2 þ 2=5B2O3  2=5TiB2 þ O2→2=5TiO2 þ 2=5B2O3 1=3ZrSi2 þ O2→1=3ZrO2 þ 2=3SiO2  ZrO2 þ SiO2→ZrSiO4  ð5Þ  ð6Þ  ð7Þ  ð8Þ  Opeka et al. [47] have reported that SiC-containing ZrB2 ceramics had high oxidation resistance up to 1500 °C compared to pure ZrB2 ceramics. Monteverde et al. [27] have reported that of ZrB2 + 5% Ni composites get degraded on oxidation at 1000 °C due to fast oxidation at the Ni rich grain boundaries. Monteverde et al. [48] have reported that oxidation of monolithic ZrB2 starts at 420 °C. It is observed that the introduction of SiC particles markedly improves oxidation resistance due to the formation of an adherent and protective borosilicate glass layer that coats the sample surface, effectively limiting the inward diffusion of oxygen toward the reaction interface. Guo et al. [49] have reported that oxidation of ZrB2 powder follows para-linear kinetics in air at 650-800 °C, where the dominating term is the parabolic one, accounting for oxygen diffusion in the oxide scale.  Fig. 15. SEM images of the oxidized (900 °C, 64 h) surface of ZrB2 with (a) 10%TiSi2, (b) 10%TiSi2 + 10%HfB2, and (c) 10%TiSi2 + 20%HfB2.  Guo et al. [50] have observed that addition of Si improves the oxidation resistance of ZrB2 at 1500 °C whereas Zr addition decreases the oxidation resistance. Karlsdottir et al. [46] have reported the formation of ZrO2 islands in the borosilicate glass on oxidation of ZrB2 + SiC ceramic at 1500 °C.  4. Conclusion  Charge composition and synthesis parameters were optimized to obtain pure ZrB2 powder by boron carbide reduction of ZrO2. Monolithic  \\x0c', 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  29  Fig. 16. Typical EDS pattern of the oxidized surface in ZrB2 + 10% TiSi2 + 20% HfB2 composites.  [5]  Fig. 17. XRD pattern of the surface oxidized at 900 °C for 64 h in air.  ZrB2 was hot pressed to full density at 1850 °C and 35 MPa. Addition of 10 wt.% TiSi2 to ZrB2 resulted in an equally high density of 98.9% by hot pressing at a lower temperature (1650 °C) and pressure (20 MPa). Similar densities were obtained for ZrB2 + HfB2 mixtures also. Hardness of monolithic sample is 23.95 GPa which decreases to 19.45 GPa on addition of 10% TiSi2. On further addition of 10% HfB2, the hardness increased to 23.08 GPa, which is very close to that of monolithic ZrB2. Fracture toughness of monolithic sample is 3.31 MPa m1/2, which increases to 6.36 MPa m1/2 on addition of 10% TiSi2. Further addition of HfB2 marginally increased the value of KIC to 6.44 and 6.59 MPa m1/2 with 10% and 20% HfB2 addition respectively. Fracture mode was found to be intergranular. Crack propagation in composites has shown considerable deflection indicating high fracture toughness. TiSi2 has been found to be a useful additive in sintering ZrB2. This helps in reducing the hot pressing temperature by 200 °C. This composite and also the addition of HfB2 have been found to have a better oxidation resistance at 900 °C. The glassy phase on the oxidized surface has been found to be a mixture of silica and silicates.  References  [1]  Fahrenholtz WG, Hilmas GE. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90(5):1347-64. [2] Bauccio ML. ASM engineered materials reference book. United States of America: ASM International; 1994. 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},{
  "_id": 115,
  "PDF": "Kinetics and evolution of oxide scale during various stages of isothermal oxidation at 1300 C in spark plasma sintered ZrB2 – SiC – LaB6 composites.pdf",
  "Text": "['Journal of the European Ceramic Society 40 (2020) 4997-5011  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e lsev ie r .com / loca te / jeu rce ramso c  Original Article  Kinetics and evolution of oxide scale during various stages of isothermal oxidation at 1300 °C in spark plasma sintered ZrB2 - SiC - LaB6 composites  T  Sunil Kumar Kashyapa,*, Ankit Kumarb, Rahul Mitraa  a Department of Metallurgical and Materials Engineering, b Department of Metallurgical and Materials Engineering,  Indian Institute of Technology Kharagpur, Kharagpur 721302, West Bengal, India Indian Institute of Technology Roorkee, Roorkee 247667, Uttarakhand, India  A R T I C L E  I N F O  Keywords: ZrB2-SiC-LaB6 composites Ultra-high temperature ceramics Spark plasma sintering Isothermal oxidation Kinetics  A B S T R A C T  The kinetics and oxide scale evolution during isothermal exposure of spark-plasma sintered ZrB2-20 SiC-LaB6 (7, 10 or 14 vol%) composites at 1300 °C for 1, 8, or 24 h have been examined. Random mass change observed during the first hour stems from a non-protective scale. The variation of mass gain with time is expressed by near-parabolic rate law during 0−8 h time period, and by relations indicating slower kinetics during 8−24 h, with parabolic rate constants (kp) decreasing sharply. Microstructural examination has shown a continuously evolving layered oxide scale comprising La2Si2O7, borosilicate glass (BSG), ZrSiO4, and ZrO2, where the BSG layer growth rate scales linearly with LaB6 content during 0−8 h, but shows an opposite trend on further exposure. During both time periods (0−8 and 8−24 h), kp decreased following a linear relationship with increasing BSG layer thickness, indicating its key role as diffusion barrier for oxygen.  1.  Introduction  Ultra-high temperature ceramics (UHTCs) developed in the 1960s, are known for their application in thermal insulation tiles or thermal barrier coatings, because of their high melting points, high thermal conductivity, high ablation resistance, and chemical attack resistance [1,2]. Zirconium diboride is one of the materials, which has been extensively considered for use as a UHTC [3]. Among the various UHTCs including borides, nitrides, and carbides, the use of ZrB2 is preferred due to its moderate density (6.1 g/cc) and high melting temperature (3250 °C), which favors the application of these materials in leadingedges of re-entry vehicles and rocket nozzles [4,5]. As reported in the literature, it is very difficult to densify monolithic zirconium diboride due to strong covalent bonding, and therefore to deal with such a problem, various types of reinforcements or additives have been used to promote densification and also lower the process temperature [6,7]. Generally, carbides (SiC, B4C), nitrides (AlN, Si3N4), and di-silicides (MoSi2, TaSi2) have been used as sintering aids for improving the densification of ZrB2 [8-13]. For the UHTCs, to sustain their structural integrity during the functioning, high oxidation resistance is considered as a pivotal requirement. Earlier studies have confirmed that the ZrB2 -SiC composites possess high oxidation resistance due to the formation of a protective outer layer of borosilicate glass (BSG) in the temperature range of 1000 °C-1700 °C. With the formation of an outer BSG layer, an  inner SiC-depleted region below the intermediate ZrO2-SiO2 layer is formed inside the oxide scale of the composite, as expected on considering the thermodynamic stability of the oxidation products [14-16]. The improved oxidation resistance is facilitated by the formation of an adherant and compact oxide scale enriched in BSG, which acts as an effective barrier against the inward diffusion of oxygen [17,18]. Along with improving the oxidation resistance, the presence of SiC particles also resists the grain growth of the ZrB2 phase during sintering, with benefits of superior thermo-mechanical properties as compared to the boride alone [19]. Some of the earlier works on the ZrB2 based composites have shown that the isothermal oxidation tests carried out for reasonably long durations exhibit mass gain kinetics following a parabolic rate law at temperatures < 1100 °C; whereas deviation towards a near-linear relationship has been observed for oxidation kinetics at higher temperatures [15,20-22]. However, the addition of SiC is found to change the mass gain kinetics from liner to parabolic in the temperature range of 1300−1600 °C by providing an effective barrier for oxygen diffusion through the oxide scale [14,23]. Liu et al. [24] have reported that the oxidation kinetics of ZrB2-20 vol% SiC composite follow a parabolic rate law; whereas on addition of 20 vol% ZrC at the expense of ZrB2, it shows near-linear kinetics at 1600 °C. Furthermore, the ZrB2-30 vol% SiC composite has followed the cubic rate law at 1500 °C due to increased viscosity along with reduction in the amount of ZrO2 inclusions  ⁎ Corresponding author. E-mail address: sunil.kashyap90@gmail.com (S.K. Kashyap).  https://doi.org/10.1016/j.jeurceramsoc.2020.07.053 Received 7 June 2020; Received in revised form 17 July 2020; Accepted 20 July 2020 Available online 25 July 2020 0955-2219/ © 2020 Elsevier Ltd. All rights reserved.  \\x0c', \"(4.0 ± 1.4 μm), LaB6 (7.9 ± 2.8 μm), and B4C (1.84 ± 1.3 μm) were used to prepare the composites for this study. The purity of all the powders was ≥99.5 %. The powders of ZrB2 and SiC were obtained from the H.C. Starck, GmbH, Goslar, Germany, whereas the LaB6 and B4C powders were purchased from the Alfa Aesar, Massachusetts, USA, and Boron Carbide India Limited, Mumbai, India, respectively. Furthermore, phenolic resin (phenol-formaldehyde with ethanol as solvent), type ABRON PR100 obtained from ABR Organics Limited, Hyderabad, India, has been used as a source of carbon in the raw materials. The carbon content of the phenolic resin has been found to be approximately 37 wt% by thermogravimetric analysis. Three different composite powder mixtures (ZSBCL-7, ZSBCL-10, ZSBCL-14) having compositions of ZrB2 + 20 vol % SiC + 5.6 vol % B4C + 4.8 vol % Carbon, with varying vol% of LaB6 (7%, 10 % and 14 %) were prepared. These powders were mixed using a planetary mono-mill (model Pulverisette 6, Fritsch GmbH, Idar-Oberstein, Germany) with acetone as the medium using vials and balls of WC-Co. After the completion of milling for 6 h at 250 rpm, the acetone was drained out, and the blended powders were dried at 350 °C for 2 h in the furnace. The composite powders were consolidated into pellets of 20 mm diameter by spark plasma sintering (Dr. Sinter, Japan) for 7 min at 1600 °C with application of ram pressure equal to 50 MPa. The bulk density of the sintered composites was measured by Archimedes' principle, whereas the theoretical densities were calculated by the rule of mixture method. Samples with dimensions of 6 mm x 4 mm x 4 mm were cut from the pellets of spark plasma sintered (SPS) composite using EDM wire-cutting, and then polished on emery paper, followed by cloth polishing with 6 μm diamond paste. Each sample was ultrasonically cleaned in acetone, and the constituent phases having different compositions were identified with the help of X-ray diffraction using a Cu target (wavelength of Cu Kα radiation = 0.15418 nm) on a high resolution diffractometer (Bruker D8 Discover, Germany) operated at 0.02˚/s step size. Microstructures of all the composites were examined by a field emission scanning electron microscope (Merlin FESEM, Carl Zeiss, Oberkochen, Germany).  2.2. Oxidation tests  Isothermal oxidation tests were performed on the composites with aforementioned compositions for 1 h, 8 h, and 24 h at 1300 °C using an electric vertical chamber furnace (designed and manufactured by BYSAKH & COMPANY, Kolkata, India) with arrangement of resistive heating by SiC elements. The change in mass during isothermal exposure was measured at selected intervals by a mechanical balance  in BSG, as reported in an earlier study [23]. The effect of the addition of La2O3 on ZrB2-20 vol% SiC composites has been investigated by Zapata-Solvas in the temperature range of 1400−1600 °C [25]. In this study, it has been found that in this temperature range, the ZrB2-20 vol% SiC composites exhibit kinetics of mass gain following a power law (Δwn = kt), where 1 ≤ n ≤ 2, indicating oxidation rates to be intermediate between those expected from linear and parabolic kinetics; whereas the addition of La2O3 to the composite has led to the kinetics based on the parabolic rate law, indicating a significant increase in oxidation resistance within the same temperature range [25]. Interestingly at 1600 °C, the oxidation exponent has been found to be 8, which is attributed to the formation of new protective (Zr/Si)OxCy phase. The addition of LaB6 to the ZrB2-SiC composites is known to contribute to the oxidation resistance by forming a refractory La2Zr2O7 and/or La2Si2O7 scale on the outermost surface of oxide scale [6,17,26,27]. Moreover, the increase in B-content due to LaB6 addition is expected to contribute an additional amount of B2O3, which in turn would enhance the fluidity and consequently, the self-healing capacity of the BSG-rich scale. Zhang et al. [26] have performed ablation tests at 2400 °C by using an oxyacetylene torch and reported that the addition of 10 vol% of LaB6 to the ZrB2-SiC composites remarkably enhanced the oxidation resistance by forming a compact La2Zr2O7 scale. Further, the oxidation resistance of the same composition was investigated by Jayaseelan et al. [27] at 1600 °C for 1 h in air. They have reported that the oxygen diffusivity through the La2Zr2O7 scale is lower than that of ZrO2, but it is higher as compared to that in the BSG. Both high viscosity and high melting temperature of the La2Zr2O7 oxide formed on the LaB6 containing composites during high temperature exposure, can be considered as the source of motivation to consider these composites as the suitable candidate material for the hypersonic applications. On the other hand, Monteverde et al. [17] have reported the absence of La2Zr2O7 on the outermost surface of the oxide scale during the arc-jet test of ZrB2-SiC-LaB6 composites at 1700 °C. The lack of La2Zr2O7 has been attributed to the high heat flux and a short duration of exposure, which suppresses the zirconate formation, thereby reducing the oxidation resistance. Further, it has been reported that the addition of LaB6 is detrimental, since it reduces the eutectic temperature of the oxide scale, and increases the oxygen vacancy concentration [6]. Although the oxidation behavior of ZrB2-SiC composites has received quite extensive attention, information about the effect of rare earth compound addition on the kinetics of oxidation is scarce [25]. Moreover, most of the earlier studies on the addition of LaB6 to the ZrB2-SiC composites have been investigated for exposure temperatures ≥1600 °C, none of those have focused on the kinetics of oxidation. Moreover, very limited understanding is available about early and intermediate stages of oxidation, which are known to control the later stages of oxide scale formation. Furthermore, the mass change kinetics and its relationship with the oxide scale evolution during initial and intermediate durations of exposure at a given temperature also needs to be examined for understanding how the mechanisms of protection against oxidation evolve with time. In view of the aforementioned aspects, this paper reports the results of investigation on the effects of varied LaB6 additions at various stages of oxidation including early, intermediate, and after exposure for 24 h, with emphasis on the evolution of the oxide scale at each stage. For this purpose, the ZrB2-20 vol % SiC composites processed by spark plasma sintering with B4C and C as additives, and having 7, 10 or 14 vol% LaB6 have been isothermally exposed for 1 h, 8 h, or 24 h at 1300 °C in dry air, and the evolution of oxide scale has been investigated.  2. Experimental Procedure  2.1. Preparation and characterization of composite  Commercially  available  powders  of  ZrB2  (5.4 ± 2.2 μm),  SiC  \\x0c\", 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 1. SEM images showing the microstructure of spark plasma sintered (a) ZSBCL-7, (b) ZSBCL-10, and (c) ZSBCL-14 composites; as well as (d) the presence of LaBO3 in ZSBCL-14 composite along with (e) EDS point analysis spectra showing the peaks from the constituent elements present in all the phases.  microstructures and compositions was carried out by FESEM and EDS. As a part of this effort, the cross-sections of the oxide scales were examined by FESEM with EDS to identify the composition beneath the outer surface till the oxide-composite interface.  3. Results  3.1. Microstructural observation  The density of ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites was measured by Archimedes principle, and the corresponding values of relative density (ratio of experimentally measured density to the theoretical density calculated on the basis of rule of mixtures) were found to be 98 %, 99 %, and 99.5 %, respectively, as reported in the author’s previous work [28]. Furthermore, the XRD patterns of the composites sintered by SPS at 1600 °C with 50 MPa pressure have shown the presence of ZrB2, SiC, and LaBO3 phases, which were found to be similar to those sintered at 1800 °C with 50 MPa pressure, as discussed in an earlier publication [29]. The SEM images of the metallographically polished surfaces depicting the microstructures of the ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites as shown in Fig. 1(a)-(c), respectively, exhibit a reasonable uniformity in the distribution of ZrB2, SiC, and B4C phases in the composites. Furthermore, the presence of LaBO3 is visible in the magnified image of the ZSBCL-10 composite, as shown in Fig. 1(d). The presence of these phases has been qualitatively confirmed by EDS elemental analysis, as shown in Fig. 1(e). The average ZrB2 grain sizes in ZSBCL-7, ZSBCL-10 and ZSBCL-14 composites processed by SPS at 1600 °C have been found to be 2.6 ± 0.6 μm, 3.7 ± 0.5 μm, and 2.7 ± 0.6 μm, respectively, as reported in a previous study [28]. The amount of SiC in the microstructures of the sintered composites [Fig. 1(a-c)] is found to be more or less unchanged at 20 vol% irrespective of the LaB6 content, as confirmed through quantification by  image analysis with the help of Image-J software. Although it was possible to identify the B4C occasionally in the microstructure by EDS analysis, yet its peaks could not be detected in the XRD patterns from the sintered composites, because its amount was low (< 5 vol%) due to its partial consumption for reduction of the oxide impurities [31]. Moreover, the variation between the distributions of SiC and B4C in the investigated composites with different amounts of LaB6 has been found to be quite insignificant. Small peaks of LaBO3 could be observed in the XRD patterns from the as-sintered ZSBCL-7, ZSBCL-10 and ZSBCL-14 composites, with their intensities scaling with the LaB6 content, as has been reported in an earlier study by the authors [29]. Formation of LaBO3 in the microstructure has been ascribed to a part of LaB6 having reacted with the surface oxides of powders. Such scavenging of oxygen is found to have a significant role in the process of densification during sintering [28,29]. Consumption of LaB6 to form LaBO3 may be considered to have lowered the amount of former compound in the microstructures of the sintered composites. Quantitative image analysis of the SEM (BSE) images depicting the microstructures, carried out by using the Image-J software on the basis of atomic number contrast has shown the amount of LaBO3 to be in the range of 5−7 vol% in case of the investigated composites. Although B2O3 and SiO2 were present as impurities in raw materials, their amounts have been found to be significantly reduced in the sintered composites due to reduction of these oxides by B4C and C, as well as scavenging action of LaB6 [28]. Moreover, in tune with this observation, the ZrB2 matrix grain boundaries as well as ZrB2-SiC interfaces in the investigated composites have been found to be sharp and free of any glassy phase.  3.2. Oxidation behavior  3.2.1. Mass change The variation  4999  of mass  change  observed  in  the  investigated  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 2. Plots of change in mass with respect to time at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h isothermal duration.  composites with respect to time, upon isothermal exposure for 1 h, 8 h or 24 h duration at 1300 °C, are shown in Fig. 2(a)-(c), respectively. Isothermal exposure of the Ni-Chrome and platinum wires at the abovementioned temperatures have shown very negligible mass change per unit surface area compared to that observed in case of the composites [18]. The results depicted in Fig. 2(a) show a significant amount of fluctuation in the trend observed for mass change within the first hour of exposure, probably because a protective scale is not able to develop during the early stage of oxidation. Interestingly, the plots in Fig. 2(b) indicate a higher mass gain for the ZSBCL-14 than that recorded for the ZSBCL-10 composite, whereas those in Fig. 2(c) indicate an opposite trend with respect to that observed after exposure for 8 h. The results in Fig. 2(c) show that with increase in the duration of exposure up to 24 h, the net mass gain experienced by the ZSBCL-14 composite is less with greater flattening of the curve, when compared with that of the ZSBCL10. However, it is clear from the graphs in both Fig. 2(b) and (c) that the ZSBCL-7 composite shows a higher mass gain than the other two composites at both the oxidation temperatures for any of the investigated isothermal durations.  3.2.2. Phase identification of oxides scale by XRD analysis The X-ray diffraction (XRD) patterns obtained from the oxide scales formed on the samples of ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites exposed at 1300 °C for 1 h, 8 h, and 24 h duration, are shown in Fig. 3(a)-(c), respectively. The XRD pattern from the oxide scale formed after exposure at 1300 °C for 1 h (Fig. 3(a)) shows the peaks corresponding to ZrO2 (00-037-1484), La2O3 (00-022-0369), and La2Si2O7 (01-072-2456). On the other hand, the XRD patterns from the oxide scale formed after exposure at 1300 °C for 8 h or 24 h show an additional peak of ZrSiO4 (01-083-1374) along with all the aforementioned peaks, but with the exception of La2O3. The difference in between the peak intensities of a given phase in the XRD patterns obtained from the oxide scales formed on the composites exposed at 1300 °C for 8 h or 24 h as shown in Fig. 3(b) and (c), respectively, might be due to the changes in the relative amounts of phases present in the oxide scales, as well as variations in oxide scale thickness and surface roughness altering the depth of location of the crystalline oxide constituents.  3.2.3. Study of oxide scale by SEM 3.2.3.1. Oxide surface morphology. The SEM images depicting the top surfaces of the oxide scales formed on the ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites exposed at 1300 °C for 1 h are shown in Fig. 4(a)-(c), respectively. The magnified views of the selected locations in the aforementioned SEM images have been also incorporated as inset. It is apparent that these surfaces are rough and porous. Comparison of the low magnification SEM images in Fig. 4, depicting the oxide scales formed on the ZSBCL-14 is more compact and less porous compared to that of ZSBCL-7 or ZSBCL-10. The results of EDS analysis, as shown in Fig. 4(d), along with the XRD results from the oxide scales formed after oxidation for 1 h have shown evidence for the formation of ZrO2, ZrSiO4, and La2Si2O7. Additionally, the presence of SiO2 phase has been confirmed by EDS analysis, and this inference is further substantiated by the presence of a hump at diffraction angles (2θ) (≤35° for 8 h exposure and ≤40° for 24 h exposure) in the XRD patterns, indicating the presence of amorphous (glassy) phase. Interestingly, such a hump is not distinct in the XRD patterns from the oxide scales formed on the composites exposed at 1300 °C for 1 h, probably because the amount of amorphous phase is very low. The presence of ZrSiO4 phase in the oxide scale formed on exposure at 1300 °C in the present study is in agreement with the results reported in an earlier study [30]. The FESEM images of the samples exposed at 1300 °C for 8 h or 24 h show the formation of a dense scale comprising La2Si2O7 appearing to be clustered or forming a network in the BSG matrix, as shown in Figs. 5 and 6. In each of these figures, ZSBCL-7, ZSBCL-10, and ZSBCL-14 are represented as (a), (b), and (c), respectively. It can be inferred from these images that the formation of the La2Si2O7 phase becomes more predominant in comparison to other oxide scale constituents, as the LaB6 content of the composite is increased to ≥10 vol%. Whereas the oxide scale formed after exposure at 1300 °C for 8 h on the ZSBCL-7 has exhibited the presence of La2Si2O7 phase as islands (Fig. 5(a)), those on the composites with higher LaB6 content appear to be in the form of large clusters or networks. Furthermore, these small islands have combined during 24 h of isothermal exposure to form big clusters of La2Si2O7, whereas ZrSiO4 has formed due to reaction between SiO2 and ZrO2 at the surrounding locations, as shown in Fig. 6(a). Intuitively, the  Fig. 3. XRD patterns obtained from the oxide scales formed during isothermal oxidation tests in the air at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h duration.  5000  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 4. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 1 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites; as well as (d) EDS point analysis spectra showing the peaks from the constituent elements present in all the phases formed after oxidation.  amount of La2Si2O7 phase formed has increased with increase in the duration of isothermal exposure from 1 h to 24 h. The XRD patterns of oxide scale (Fig. 3) have also shown the intensity of La2Si2O7 peaks to scale with both isothermal test duration, and LaB6 content of the composite. The absence of peaks belonging to La2O3 in the XRD patterns from the oxide scales formed by exposure for 8 h or 24 h at 1300 °C is ascribed to their consumption in the reaction with SiO2 to form La2Si2O7.  3.2.3.2. Cross-section of oxide scale. A typical SEM (BSE) image depicting the cross-section of the oxide scale formed on the ZSBCL-14 composite due to the exposure at 1300 °C for 1 h is shown in Fig. 7(a), whereas the EDS X-ray maps depicting the locations of enrichment of Zr, Si, O, La, and B are shown in Fig. 7(b-f), respectively. Examination of the aforementioned EDS X-ray maps as shown in Fig. 7(b-f) indicates the formation of an oxide scale comprising ZrO2-B2O3-SiO2 mixture.  Typical backscattered electron (BSE) micrographs depicting the crosssections of the oxide scale formed by exposure at 1300 °C for 1 h on the ZSBCL-14 at higher magnification indicates that the oxide scale comprises a SiO2-rich outer layer of BSG with inclusions of ZrO2 having a greater area fraction near the interface of this layer with a porous inner layer composed of a ZrO2+SiO2-B2O3 mixture. It is interesting to note that no La-containing oxide is visible as an outer layer in the cross-section image of the oxide scale formed after 1 h duration, and this could be because of its presence as discrete islands with varied sizes, and not as a continuous film at the top surface depicted in Fig. 4(c). On examination of the SEM images in Fig. 8(b) and (c) showing the oxide scales formed on exposure at 1300 °C for 8 h and 24 h, respectively, it is evident that the outer layer is made of a thin but more or less continuous La2Si2O7, followed by inner layers comprising the following oxides sequentially arranged: BSG, ZrO2, and porous ZrO2 with traces of BSG. Further, the thickness of BSG  Fig. 5. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 8 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites.  5001  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 6. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 24 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites.  layer appears to have increased with increasing duration of the isothermal hold at 1300 °C. Based on the comparison of the SEM images depicting the cross-sections of the oxide scales formed at 1300 °C for durations from 1 to 24 h, it is appropriate to infer that the growth of a distinct protective outer layer of La2Si2O7 over the composites is possible only after a minimum duration, which may be related to their LaB6 content. The thicknesses of the oxide scales formed on the investigated composites after exposure for 1 h, 8 h, and 24 h at 1300 °C are presented in Table 1. Based on the results shown in this table, the following inferences may be drawn: (i) on exposure for 1 h or 8 h, the thickness of the oxide scale formed on the ZSBCL-14 is the highest, followed by those on ZSBCL-7 and ZSBCL-10; and (ii) thicknesses of the ZrO2 layer in the oxide scales of the investigated composites are found to be comparable on exposure for 8 h or 24 h; whereas (iii) thickness of BSG layer in the oxide scales increases in the following order ZSBCL7 < ZSBCL10 < ZSBCL14. In other words, after 24 h of exposure, the ZSBCL-14 exhibited the lowest oxide scale thickness with maximum BSG formation amongst the investigated composites. The lowest total oxide scale thickness observed in case of the ZSBCL-10 after 1 or 8 h, and of ZSBCL-14 after 24 h of isothermal exposure appear to be consistent with their lowest mass gain among the investigated composites during isothermal hold for the respective durations as shown in Fig. 2.  4. Discussion  4.1. Thermodynamic parameters of oxidation reactions  The thermodynamic aspects of the oxide scale formation during isothermal exposure at 1300 °C have been analyzed. The values of change in free energy of formation per mole of oxygen [ΔGf (kJ/mol)], the equilibrium constant (Kf,eq), and the partial pressure of oxygen ( ) have been calculated with the help of data available in the literature [31-33]; and presented in Table 2. The values of Kf,eq have been calculated by using the following relation:  ΔGf = -RT ln (Kf,eq)  (1)  where R is the universal temperature. Further, the values of for the formation of oxides have been calculated by substituting the values of Kf,eq in the following equations:  constant,  and T is  the  absolute  ideal  gas  (2)  (3)  Fig. 7. Cross-section of oxide scale formed on ZSBCL14 by exposure for 1 h at 1300 °C: (a) SEM (BSE) image, and EDS X-ray maps showing enrichment of (b) Zr, (c) Si, (d) O, (e) La, and (f) B.  5002  pO2pO2=()()()Kaaap**()feqZrOZrOBOZrBO,()2/52/52/5222322=()Kapap*()()*()feqSiOSiOCOSiCO,()2/32/32/3222\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 8. Cross-section of oxide scale formed on ZSBCL14 by exposure at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h duration.  (4)  (5)  where is the activity of various reactants and products mentioned in the considered reactions as outlined in Table 2, and is the equilibrium partial pressure of oxygen for the given reactions. The values of have been calculated by taking the activity of pure substance as unity, and ignoring the presence of CO in Eq. (5) [15]. The results in Table 2 indicate that all the reactions are feasible at 1300 °C, because the ΔGf (kJ/mol) values are negative. Based on the ΔGf (kJ/mol) values shown in Table 2, the chemical stability of oxides at investigated temperatures is found to decrease in the following order: La2Si2O7 > ZrSiO4 > ZrO2 > SiO2 > SiO > La2O3. The evidence for the presence of ZrSiO4 in the oxide scales as depicted in Figs. 3, 4, and 6 may be considered to be consistent with its co-existence in thermodynamic equilibrium with ZrO2, as predicted by the ZrO2-SiO2 phase diagram [34]. The absence of La2Zr2O7 may be attributed to the significantly lower (more negative) ΔGf (kJ/mol) value of La2Si2O7 (-2809 KJ/mol) as compared to that of the former oxide (-150 KJ/mol) at 1300 °C [35]. The lower free energy of formation of La2Si2O7 may have preferably promoted its formation, rather than that of La2Zr2O7 during exposure at 1300 °C. It may be noted that a few earlier studies [6,17,26,27] have reported the presence of the La2Zr2O7 in the oxide scales of ZrB2-SiC-LaB6 composite exposed only at temperatures ≥1600 °C, and therefore its absence from the oxide scales formed at 1300 °C in the present study is not surprising. Based on the data presented in Table 2, the free energy of the Reaction (2) is more negative than the Reaction (4), which indicates that the formation of SiO2 is more strongly favored compared to that of SiO at the investigated temperatures. The formation of SiO at temperatures > 1500 °C by active oxidation of SiC causing the formation of a SiC-depleted region, at the oxygen-deficient locations (or regions having lower partial pressure of oxygen) underneath the top surface, has been reported by many researchers [15,36-38]. However, a few authors have reported that the SiC-depleted region is formed even at a  of oxygen  low partial pressure  lower temperature 1200 °C under [39,40]. A high value of for La2O3 formation indicates a relatively lower driving force for the oxidation reactions leading to its formation, which can be easily correlated with limited growth of the La-based oxide scale (as shown in Fig. 8) during oxidation. It has been shown in earlier reports that the oxidation of LaB6 starts at 700−800 °C, resulting in the formation of a few intermediate oxides, and those oxides may convert to La2O3 at temperature ≥1200 °C [21,41]. On the other hand, the oxidation of ZrB2 and SiC starts at ≥800 °C and ≥1100 °C, respectively [15].  4.2. Oxidation kinetics  The oxidation kinetics of the composites have been analyzed on the basis of the data representing the change in mass (ΔW) by oxidation during isothermal exposure (t) to calculate the oxidation rate-related parameters such as the oxidation exponent (n), the general rate constant (k), and the parabolic rate constant (kp). The oxidation kinetics can be determined by using the power-law equation:  (ΔW)n = kt  Taking logarithm of both the sides,  log (ΔW) = (1/n).  log k + (1/n).  log t  (6)  (7)  The value of n can be obtained by the best-fit line of the slope for the log-log plots of (ΔW) against t. Further, the value of the general rate constant (k) can be determined by the intercept of the aforementioned best fit line to the y-axis. If the mass gain within a particular interval of exposure follows parabolic rate law, then the value of n in Eq. (6) would be 2, whereas a typical linear relationship would lead to a value of n1 [42,43]. The parabolic rate law implies that the oxidation reactions are controlled by diffusion of oxygen (atomic or anionic oxygen) from oxide-air interface to the oxide-composite interface through the protective oxide layer [44,45]. Assuming that the parabolic law has been followed in the current investigation, the value of parabolic rate constant (kp) can be determined by the following relation:  Table 1 Oxide scale thickness of composites after isothermal hold for 1 h, 8 h, and 24 h at 1300 °C.  Composition  Thickness of oxide scale formed after 1 h of exposure (μm) BSG + ZrO2  ZSBCL-7 ZSBCL-10 ZSBCL-14  57 ± 9 45 ± 2.3 86 ± 8  Thickness of oxide scale formed after 8 h of exposure (μm)  Thickness of oxide scale formed after 24 h of exposure (μm)  BSG  ZrO2  Porous ZrO2 + BSG  Total  BSG  ZrO2  Porous ZrO2 + BSG  Total  12 ± 2 17 ± 2.2 27 ± 4.3  11 ± 1 10 ± 1 12 ± 2.7  64 ± 3 55 ± 1.6 53 ± 3.2  87 ± 3.7 82 ± 2.9 92 ± 6  23 ± 4.3 26 ± 3 31 ± 1.7  13 ± 1.3 14 ± 1.5 13 ± 1.4  93 ± 3.7 73 ± 1.3 59 ± 6.5  129 ± 5.8 113 ± 3.6 103 ± 6.8  5003  =()()()Kaaap**()feqLaOLaOBOLaBO,()2/214/74/2123232362=Kppap()*()()*()feqSiOSiOCOSiCO,()2apO2pO2pO2\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Table 2 The values of change in free energy per mole of oxygen (ΔGf), the equilibrium constant (Kf,eq), and the partial pressure of oxygen ( 1300 °C. Free energy of formation of oxidation products, ZrSiO4 and La2Si2O7 are also included.  ) for the oxidation reactions at  NO.  1 2 3 4  5 6  Reactions  (2/5)ZrB2 (s)+ O2 (g) = (2/5)ZrO2 (s) + (2/5)B2O3 (l) (2/3)SiC (s) + O2 (g) = (2/3)SiO2 (l) + (2/3)CO (g) (4/21)LaB6 (s) + O2 (g) = (2/21)La2O3 (s) + (4/7)B2O3 (l) SiC (s) + O2 (g) = SiO (g) + CO (g) Free energy of formation for ternary oxidation products. La2O3 + 2SiO2= La2Si2O7 ZrO2 + SiO2 = ZrSiO4  ΔGf (KJ/mol)  −566.9 −548.6 −132.9 −418.7  −2809.0 −1432.6  Kf,eq  6.7 × 1018 1.6 × 1018 2.6 × 104 8.0 × 1013  (atm)  1.5 × 10−19 6.0 × 10−19 3.8 × 10−5 1.2 × 10−14  Fig. 9. Typical plots of composites oxidized at 1300 °C showing the logarithmic plots of mass change (ΔW) against time for (a) 8 h and (b) 24 h; as well as plots of ΔW2 against time for (c) 8 h and (d) 24 h duration.  (ΔW)2 = kp.t + c  (8)  where kp is the parabolic rate constant, and c is a constant. Further,  d(ΔW)/dt = kp/2(ΔW)  (9)  The value of kp can be determined by the best-fit line of the slope for the plots of ΔW2 against t. The value of kp (= square of mass gain at unit time) is known to scale with the rate of mass gain by oxidation. Typical plots of log (ΔW) against log (t) for the isothermal duration of 8 h and 24 h are shown in Fig. 9(a) and (b), respectively, whereas those showing the variation of (ΔW)2 against t for the aforementioned duration are shown in Fig. 9(c) and (d). The value of oxidation exponent, general rate constant, and parabolic rate constant of the composites exposed at 1300 °C for 8 h and 24 h are presented in Tables 3 and 4, respectively. From the results in Table 3, the following inferences may be drawn: (i) the value of the general rate constant obtained for the ZSBCL-7 is an order of magnitude greater than that of either ZSBCL-10 or ZSBCL-14  Table 3 The values of oxidation exponent, the general rate constant, and the parabolic rate constant of composites oxidized at 1300 °C for 8 h.  Composite  ZSBCL-7 ZSBCL-10 ZSBCL-14  Time span (h)  1−8 1−8 1−5 5−8  Oxidation exponent (n)  General rate constant (k) (mgn cm−2n h-1)  2.8 2.1 1.9 3.4  60.56 8.1 8.2 118.7  R2  .98 .99 .99 .90  Parabolic rate constant (kp) (mg2 cm−4 h-1)  8.8 6.4 9.7 4.8  R2  .98 .99 .99 .84  composite; and (ii) the rate of oxidation of the ZSBCL-14 composite becomes slower (n = 3.4) after 5 h of isothermal exposure, which is consistent with the change in slope of the corresponding best-fit line in Fig. 9(a). Furthermore, considering that the values of n to be close to 2 for the initial stage of exposure (i.e. 8 h) for ZSBCL-7 or ZSBCL-10, and 5 h for the ZSBCL-14, the approximation involved in the use of  5004  pO2pO2\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Table 4 The values of oxidation exponent, the general rate constant, and the parabolic rate constant of composites oxidized at 1300 °C for 24 h.  Composite  Time span (h)  Oxidation exponent (n)  General rate constant (k) (mgn cm−2n h-1)  ZSBCL-7  ZSBCL-10  ZSBCL-14  1−24  1−24  1−8 8−24  3.9  3.0  2.9 9  486  45  45.6 9.9 × 106  R2  .99  .99  .99 .98  Time span (h)  Parabolic rate constant (kp) (mg2 cm−4 h-1)  1−8 8−24 1−8 8−24 1−8 8−24  7 3 5.4 2.8 5.6 0.9  R2  .99 .99 .99 .99 .97 .97  parabolic rate law as depicted in Fig. 9(c) and (d) appears quite appropriate. The results presented in Fig. 9(c) and listed in Table 3 show the value of kp as 9.7 mg2 cm−4 h-1 for the ZSBCL-14, which indeed is the highest among the investigated composites exposed at 1300 °C for 8 h, but is reduced to 5.5 mg2 cm−4 h-1 for the later stages of exposure, which is less than the values of kp for ZSBCL-7 (8.8 mg2 cm−4 h-1) and ZSBCL-10 (6.4 mg2 cm−4 h-1). Such a trend may also be considered as qualitatively consistent with the results of the mass gain against time (Fig. 2(c)). The values of n in the range of 3−4 for all three investigated composites exposed for 24 h at 1300 °C, as shown in Fig. 9(b) and Table 4, indicate that the average rate of oxidation during exposure for 24 h is significantly slower than that predicted by the parabolic rate law. Furthermore, the highest value of n (9) observed for the ZSBCL14 during the 8−24 h span of isothermal exposure indicates that its rate of mass change is significantly slower than that of ZSBCL-7 or ZSBCL-10, beyond the initial time-period of 8 h. It may be further noted that the plots in Fig. 9(d) represent the two different regimes of the rate of oxidation. In the first regime, which represents the early stage of oxidation (up to 8 h), the rate of oxidation of all the investigated composites is higher compared to that in the second regime (8-24 h). The higher rate of oxidation in the early stage of oxidation (≤8 h) may be attributed to the formation of an inadequately protective oxide scale during the initial stage of oxidation. Furthermore, in the second regime (8−24 h), the value of kp has been found to be similar for ZSBCL-7 and ZSBCL-10, whereas a significantly lower value is obtained for the ZSBCL-14 composite. Plots of the rate of mass change (dW/dt) calculated from Eq. (9) against time of isothermal exposure as shown in Fig. 10, indicate that (i) the rate of mass gain for the ZSBCL-14 composite is lower than that of ZSBCL-7 or ZSBCL-10 after the initial duration (9 h) of exposure, and the observed trend is similar to that observed for the kinetic parameters such as n and kp; (ii) for each of the investigated composites, a steady state characterized by a sharp decrease in the value of d(ΔW)/dt is reached beyond a limiting duration of 10−12 h; (iii) the transition in the slope of d(ΔW)/dt vs time plots is reached relatively earlier and the steady-state condition is most welldefined in the ZSBCL-14 composite; and (iv) at the end of 24 h, the values of d(ΔW)/dt decrease in the following order, ZSBCL-7 > ZSBCL10 > ZSBCL-14. The commonly observed decrease in the rate of oxidation with the increase in the duration of exposure from 8 h to 24 h at 1300 °C for each of the investigated composites may be attributed to the progressive evolution of a compact and more protective oxide scale, which acts as an effective diffusion barrier for oxygen. According to the available literature, the value of n is reported to be 2 for SiC containing UHTCs at an intermediate temperature range (1300−1600 °C), and this has been rationalized by considering the diffusion of oxygen through the viscous BSG layer as the rate-limiting step [16,46]. However, it deviates from the parabolic rate law to a higher-order, if the rate of diffusion of oxygen through the protective oxide scale is further reduced [47]. Guo et al. [23] have reported that increasing the amount of SiC from 10 to 30 vol% in the ZrB2-SiC composite, changes the oxidation kinetics from parabolic to cubic due to the formation of a higher amount of SiO2 in the oxide scale. In the current study, the growth of BSG layer also plays a similar important role in oxidation kinetics. After forming an appreciable amount of BSG  layer, the oxidation exponent has been found to be 9 for the ZSBCL14 composite at 1300 °C for 24 h exposure. Furthermore, the value of kp (5−10 mg2 cm−4 h-1) obtained after 8 h of exposure in the present study is in good agreement with the results reported by Kakroudi et al. [48] on the ZrB2-SiC-TaC composites, which has shown the value of kp to vary from 2.2-30.6 mg2 cm−4 h-1 with increase in the temperature of isothermal exposure from 1000 °C to 1400 °C for 10 h duration. ZapataSolvas et al. [25] have reported that the addition of La2O3 drastically enhances the oxidation resistance with n 8 at 1600 °C by forming a protective MeOxCy (Me = Zr, Hf, or Si) phase. These results are comparable with the current study showing the value of n  9 for the ZSBCL-14 composite after forming a sufficiently thick protective oxide scale at 1300 °C. Obviously, the kinetics of oxidation at 1300 °C in the present study is expected to be slower than that observed in the earlier investigation at 1600 °C [25].  4.3. Evolution of oxide scale and its relationship with kinetic parameters  evolving  4.3.1. Effect of oxidation products on structure and integrity of BSG layer The examination of the oxide scales formed on the investigated composites after isothermal exposure for various durations at 1300 °C through XRD and SEM-EDS analyses has revealed that La2Si2O7, SiO2, ZrSiO4, and ZrO2 are the main products of oxidation. The positive component of mass change observed during oxidation is ascribed to the formation of non-volatile products including ZrO2, SiO2, La2Si2O7, and ZrSiO4, whereas the negative component can be explained on the basis of the formation of gaseous products like CO2 (g), CO (g), and B2O3 (g) followed by their escape, as well as occasional spallation of a part of the oxide scale triggered by growth-related internal stresses. It is necessary to understand the effect of various constituent phases on the formation of a protective glassy scale, which may vary with the net composition. It is well-known that increase in the amount of B2O3 in the borosilicate glass increases its fluidity and plasticity with lowering of viscosity [49]. On the other hand, its vaporization causing increase in the proportion of SiO2 enhances the oxide scale viscosity. Further, earlier studies have shown that the presence of Zr2+ ions in borosilicate glass acts as a glass network modifier and increase the glass-forming ability. The SiO4 − tetrahedral network is strengthened by lowering the fraction of non-bridging oxygen atoms created by vaporization of B2O3 due to the connections established between the ions of (SiO4)and (ZrO6)2[50,51]. It has been proposed in an earlier study that as a result of increase in viscosity due to the dissolution of a significant amount of ZrO2, the molten BSG is prevented from ascending to the top surface from the ZrO2+BSG mixed layer [27]. Further, mixing of a significant amount of undissolved ZrO2 inclusions (beyond percolation threshold) with BSG may be detrimental, as the former oxide provides easier diffusion path for oxygen anions. It has been shown that addition of a small amount of La2O3 (1 mol %) to B2O3-SiO2 (binary) system causes strong phase separation due to miscibility gaps in binary SiO2-La2O3 and B2O3-La2O3 systems [52]. In a similar manner, mixtures having higher B2O3 content (40−60 mol%) with moderate SiO2 content (20−40 mol%) have exhibited separation of the La-rich phase. On the contrary, addition of a large amount of La2O3 (40−60 mol%) to B2O3-SiO2 glass leads to crystallization with  5005  \\x0c', 'F  i  g  .  0 1  .  P  l  o  t  s  d  e  p  i  c  t  i  n  g  t  h  e  v  a  r  i  a  t  i  o  n  i  n  r  a  t  e  o  f  m  a  s s  c  h  a  n  g  e  a  g  a  i  n  s  t  t  i  m  e  f  o  r  c  o  m  p  o  s  i  t  e  s  e  x  p  o  s  e  d  i  s  o  t  h  e  r  m  a  l l  y  a  t  0 0 3 1  °  C  f  o  r  4 2  h  .  S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  5006  \\x0c', 'S.K. Kashyap, et al.  de-polymerization along with an increase in the fraction of non-bridging oxygen anions, which in turn lowers the glass-transition temperature. In such cases, the La3+ ions in La2O3 play the role of chargecompensation by occupying positions close to borosilicate network and transferring non-bridging oxygen anions to SiO4 − and BO3 units [53]. On the other hand, mixtures with almost equal amounts of SiO2 (40−60 mol%) and B2O3 (30−60 mol%) with 20−30 mol% La2O3 have formed homogeneous glass, where the La3+ ions have bonded strongly to non-bridging oxygen anions and thereby acted as glassnetwork modifier, leading to increase in the glass transition temperature. It has been earlier reported that the SiO2 tetrahedron with La atoms in interstitial positions is more stable than that having those at substitutional sites, because the interstitial La breaks the nearest Si-O bond to form La-O and La-Si bonds, which is beneficial for improving the high‐temperature stability of SiO2 [54]. In another study, it has been demonstrated that the formation of La2O3 near SiC promotes its inclusion into the BSG phase, which in turn increases both its viscosity and thermal stability, and thereby lowers the oxygen diffusivity [25]. Further, phase separation with La2O3 addition to the BSG due to miscibility gap is expected to increase the viscosity and lower oxygen diffusivity [27]. Based on the calculated isothermal section of the ZrO2-SiO2-B2O3 phase diagram at 1500 °C, increase in the B2O3 content of the BSG liquid enhances its solubility for ZrO2 resulting in the formation of ZrO2-SiO2-B2O3 (BSZ) liquid. In a similar manner, La2O3 is also expected to dissolve in the BSG liquid. However, vaporization of B2O3 is expected to cause enrichment of ZrO2 and La2O3 in the oxide scale, leading to formation of ZrSiO4 and La2Si2O7, respectively [23,27], whose presence has been confirmed in the present study by both XRD analysis and EDS examination of the phases present in the oxide scales, as shown in Figs. 3-6.  4.3.2. Effect of LaB6 content on protective character of oxide scale In the present study, the microstructural study of top surface and cross-section of the oxide scale has revealed the absence of a continuous layer of oxide scale after 1 h of isothermal exposure [Figs. 4,7 and 8(a)], which is in tune with the random fluctuations in mass change from positive to negative during the initial stage of oxidation. The oxide scale formed after exposure at 1300 °C for 1 h, is expected to contain relatively less amount of SiO2 compared to that of B2O3 due to faster oxidation of both LaB6 and ZrB2 phases. Partial vaporization of B2O3 may along with spallation have contributed to randomly observed decrease in mass, as shown in Fig. 2(a). Interestingly, in the course of isothermal exposure for 1 h, the oxide scale developed on the ZSBCL-7 composite shows the presence of open pores. On the other hand, the ZSBCL-10 and ZSBCL-14 composites show relatively insignificant cracking on the surface of the oxide scale. This observation could be attributed to the formation of a greater amount of B2O3 due to the presence of much higher amount of LaB6 in both ZSBCL-10 and ZSBCL-14 composites. It may be noted that the amount of B2O3 formed by oxidation of a mole of LaB6 is expected to be 3 times more than that by a similar amount of ZrB2. The formation of a larger amount of B2O3 is expected to increase the fluidity of the BSG scale, and facilitate self-healing of cracks and porosities. However, due to the formation of a relatively smaller amount of La2O3 compared to that of B2O3 (being contributed by oxidation of both LaB6 and ZrB2), phase separation within the BSG may be responsible for identification of the former phase in the XRD patterns from the oxide scales formed through exposure for 1 h at 1300 °C (Fig. 3(a)). At the locations having a larger amount of La2O3 formed along with BSG containing nearly equivalent amounts of B2O3 and SiO2, crystallization with formation of La2Si2O7 may have happened, and therefore this phase appears to be quiet scattered in the resultant oxide scale (Fig. 4). As mentioned above, a limited dissolution of ZrO2 in the BSG during the first hour of exposure is expected to have strengthened the glass with lowering the possibility of crystallization. A few earlier reports have shown that oxidation kinetics depends to  Journal of the European Ceramic Society 40 (2020) 4997-5011  a large extent on the presence of open pores on the oxide scale surfaces, formation of cracks and spallation of oxide scale, porous structure beneath the BSG layer, and SiC-depleted region [55,56]. Therefore, it may be speculated that due to the presence of the highest amount of LaB6 in the ZSBCL-14 composite, creation of pores in the oxide scale by vaporization of B2O3 has probably enhanced the oxidation rate during the initial stages (0−5 h) of exposure at 1300 °C, compared to that observed in other two composites. On the other hand, during the time span of 8−24 h at 1300 °C, the amount of SiO2 formation appears to have increased proportionately to the other constituents of the oxide scale, and this may have accelerated the formation of a homogenous BSG layer within the oxide scale, which is considered as primarily responsible for protection against further oxidation. As the amount of La2O3 (having more mass than an equivalent amount of ZrO2 formed by oxidation of ZrB2) and B2O3 formed in the oxide scales of the investigated composites with less LaB6 content is expected to be less too, the ZSBCL-7 appears to have exhibited the highest oxidation resistance during the span of 0−8 h. Although the relative density of the ZSBCL-7 (98 %) is slightly less than that of ZSBCL-10 (99 %) and ZSBCL-14 (99.5 %), this factor too is not expected to have a significant effect on kinetics of oxide scale formation. The average sizes of ZrB2 grains in the microstructures of the investigated composites have been found to have only minor variations (2.6−3.7 μm), depending on the amount of LaB6 content, which suggests that this may not have affected the kinetics of oxide scale growth significantly. Therefore, the role of ZrB2 grain boundaries and ZrB2-SiC interfaces appears to be similar for the investigated composites, and may not have caused much difference in the oxidation behavior.  4.3.3. Effect of LaB6 content on BSG layer growth Investigation of the oxide scales developed after 8 h or 24 h of isothermal exposure shows the well-developed and distinct layers of BSG, and La2Si2O7 on the outermost surface of typically multi-stratified oxide scales [Figs. 5, 6 and 8(b-c)]. The thickness of the BSG layer formed after 8 h of isothermal hold at 1300 °C as shown in Table 1, is found to be increased from 12 μm (in ZSBCL-7) to 32 μm (in ZSBCL-14) with increasing volume fraction of LaB6, while the ZrO2 layer thickness for all three composites is found to be almost similar (10−13 μm), irrespective of the LaB6 content. Such an increase observed in the BSG content alone with the amount of ZrO2 remaining unchanged with increasing amount of LaB6 in the composites suggests that the SiC particles at sub-surface locations are preferentially oxidized. Such preferential oxidation of sub-surface located SiC particles may be ascribed to lower partial pressure of oxygen required for oxidation of SiC compared to that for ZrB2, as shown in Table 2. Also, there is a strong possibility of active oxidation of a part of SiC through Reaction (4) in Table 2, and the escaping SiO (g) is probably converted to SiO2 (s) at the oxide-air interface. The escape of SiO(g) and CO(g) from the reaction front is expected to further enhance the oxidation rate of SiC, as is normally expected according to the law of mass action. Interestingly, the formation of porous ZrO2 with traces of BSG due to the ingress of oxygen from the outer scale appears to be partially restricted with the increasing volume fraction of LaB6. Measurements of thickness of various layers within the oxide scales after exposure at 1300 °C for 24 h have shown that the porous ZrO2 + BSG is 64 ± 3 μm thick in case of the ZSBCL-7 composite, whereas it is only 53 ± 2.2 μm thick in case of the ZSBCL-14, indicating that the oxide scale is more protective in case of the latter composite. Based on the measurements carried out using the SEM images and the data presented in Table 1, plots depicting the variation of the intermediate BSG layer thickness with LaB6 content during the time-spans of 1−8 h and 8−24 h are presented in Fig. 11(a) and (b), respectively. According to the results shown in Fig. 11(a), the growth rate of the BSG layer within the oxide scale at the end of 8 h isothermal exposure is found to scale with increasing LaB6 content of the investigated composites following a linear  5007  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 11. Plots depicting the variation of (a) BSG layer growth up to 8 h with LaB6 content; and (b) BSG layer growth during 8-24 h span with LaB6 content. Best-fit lines are drawn in both the plots.  relationship, t(BSG) = 7.5x + 3.66 with R2 = 0.96. This observation may be ascribed to the increased formation of B2O3 by oxidation of LaB6, which probably has lowered the viscosity of the BSG layer during isothermal exposure at 1300 °C, and this in turn may have enhanced the diffusivity of oxygen, leading to increase in the growth rate. However, the growth rate of the BSG layer appears to decrease with increasing LaB6 content during the span of 8−24 h following a linear relationship, as evidenced by the linear-fit of data, t(BSG) = -3.5x + 15 with R2 = 0.94 in the plot shown in Fig. 11(b). For example, the thickness of BSG layer has increased by 91 % or 54 %, respectively for ZSBCL-7 or ZSBCL-10 composites on increase of the isothermal exposure duration from 8 h to 24 h, but such increment (15 %) is significantly less in ZSBCL-14 composite. Further, by comparing the results in Fig. 11(a) and (b), it is obvious that for each of the investigated composites with given LaB6 content, the observed growth rate during the span of 8−24 h is significantly less than that recorded for the duration up to 8 h. This observation also suggests that there is a reversal in the growth rate of the BSG layer, after a critical thickness is reached for a given composite. In order to explain this observation, it is necessary to consider that with increasing duration of exposure, the amount of B2O3 dissolved in the BSG layer may have decreased by vaporization, and this resulted in increased fractions of dissolved ZrO2, La2O3, and SiO2 as a natural consequence, causing a simultaneous increase in the viscosity. It is well-expected that an increase in the viscosity of the BSG layer would significantly lower the diffusivity of oxygen. Such an increase in the volume fractions of the aforementioned non-volatile oxides may therefore be considered as the possible explanation for the observation of lower rate of growth in the BSG layer during the 8−24 h span compared to that happening on exposure within 1−8 h time-period for each of the investigated composites, irrespective of the LaB6 content. A similar reason may also be proposed to explain the observed decrease in the growth of BSG layer thickness during the 8−24 h exposure span with increase in the LaB6 content of the investigated composites.  4.3.4. Effect of BSG layer growth on oxidation kinetics After a continuous BSG layer is formed, it is appropriate to infer that diffusion of oxygen through the protective BSG layer is the rate limiting factor in case of the investigated composites, because the diffusivity of oxygen through SiO2 (10−21 m2/s at 1550 °C) is slower by several orders of magnitude than that in ZrO2 (10−10 m2/s at 1500 °C) [57,58]. In order to correlate the oxidation kinetics with growth of the oxide scales of the investigated composites, the values of parabolic rate constant, kp are plotted against the BSG layer thickness recorded after exposure for 8 h and 24 h, as shown in Fig. 12(a) and (b), respectively. As shown in both Fig. 12(a) and (b), the values of parabolic rate constant, kp decreases with increasing thickness of the BSG layer, whereas a linear relationship appears to be followed during exposure up to 8 h, such that the data fall in a single best-fit line (kp = -0.25*t(BSG) + 11.36) with R2 value of 0.91. As the oxidation rate depends on the diffusion of oxygen through the BSG layer, it is intuitive that a thicker layer of BSG would lead to increase in the diffusion distance, and therefore a smaller value of kp, as shown in Fig. 12(a) and (b). The  absence of a single best-fit line for the data in Fig. 12(b) may be ascribed to the difference between the compositions of the BSG layer formed in the oxide scales of the investigated composites, which in turn is expected to have altered the oxygen diffusivities through these layers. It may be noted that in Fig. 12(a), the selected kp values represent a span of 0−8 h for ZSBCL-7 and ZSBCL-10, whereas it is obtained for the span of 5−8 h in case of the ZSBCL-14 composite. The choice of kp representing the span of 5−8 h in case of the ZSBCL-14 composite has been made on the basis of the following considerations: (i) the oxidation resistance of the investigated composites is primarily facilitated by the formation of BSG layer within the oxide scale, and (ii) the kp value calculated in case of the ZSBCL-14 composite has decreased sharply on moving from the span of 0−5 h to 5−8 h as shown in Fig. 9(c), which indicates that the BSG has grown in thickness, or its composition has changed more drastically during the latter time-span, such that the ingress of oxygen is significantly retarded. On the other hand, the kp values observed in case of the ZSBCL-7 and ZSBCL-10 composites have remained more or less constant during the span of 0−8 h, suggesting that the BSG layer growth rate along with change in viscosity through compositional changes may have taken place somewhat uniformly throughout this time-period, except for the first hour, after which the presence of this layer could not be detected through microstructural examinations (Table 1). Further, the observed decrease in the values of the rate of mass change, d(ΔW/dt) with time (as shown in Fig. 10) after the first hour of exposure may be considered as an indirect confirmation of a continuous evolution of the protective BSG layer. As inferred from the trend in Fig. 10 and the kinetic data in Tables 3 and 4, the ZSBCL-7 has exhibited the lowest rate of mass gain during the initial 8 h of exposure, which is followed by ZSBCL-10 and ZSBCL14 composites, whereas the trend is completely reversed after an exposure duration of 8−9 h. This observation may be ascribed to faster rate of oxidation of LaB6 compared to that of ZrB2 and SiC, and therefore the rate of mass gain by oxidation in the initial stages appears to scale with the LaB6 content of the investigated composites. Once the LaB6 present near surface is almost fully converted to La2O3, which probably takes 8−9 h, the oxidation process is controlled primarily by oxidation of SiC with formation of BSG and La2Si2O7 along with increase in viscosity (as discussed in Section 4.1 and earlier part of this section), thereby enhancing the protective character of the oxide scales. Considering that the mass gain by the ZSBCL-14 is more than that of the ZSBCL-10 up to the exposure duration for 8 h, it is inferred that a critical minimum duration is required for a desirable ratio of oxide scale constituents to form, so that a protective glassy layer with desirable amount of viscosity is formed within the oxide scale. This hypothesis is in line with the observations recorded from Fig. 12(a), indicating that the faster the growth rate of the BSG layer in the oxide scales, the lower is the observed value of kp from the isothermal oxidation tests on the investigated composites. This observation is justified, considering the increase in diffusion distance of oxygen with increasing growth of the BSG layer within the oxide scale. If the time-span of isothermal exposure up to 8 h at 1300 °C is considered for the ZSBCL-14 composite, a transition in the mass gain kinetics is observed after the initial period of  5008  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  Fig. 12. Plots showing the change in the values of kp with thickness of BSG layer after isothermal exposure of (a) 8 h, and (c) 24 h duration.  5 h, as is obvious from the results shown in Table 3. Furthermore, on crossing over from the time span of 0−8 h to 8-24 h, the value of oxidation exponent (n) calculated for the ZSBCL-14 composite is increased very sharply from 2.9 to 9, whereas the parabolic rate constant (kp) is decreased from 5.6 to 0.9 mg2 cm−4 h-1. Such an observation regarding sharp changes in oxidation rate (kinetics-related) parameters is consistent with the trend showing lower rate of mass gain in the ZSBCL-14 compared to that in ZSBCL-7 or ZSBC-10 (Fig. 10), which in turn appears to be related to significantly reduced rate of the BSG layer growth in the oxide scale of the ZSBCL-14 during the span of 8−24 h. Such an observation may be explained by the requirement for the formation of a critical thickness of dense and continuous BSG layer for inhibiting the oxygen penetration and the consequent enhancement of the oxidation resistance of the composites. Obviously, the present observations suggest that such a critical BSG layer thickness is reached earliest in the oxide scale of the ZSBCL-14 composite, which in turn leads to a sharp decrease in kp after exposure for 5 h at 1300 °C.  and B2O3 formed by its oxidation, along with the continuous change in the amount of the latter constituent due to vaporization during exposure at 1300 °C. The results of the present study have significantly improved the scientific understanding of the stages of oxide scale evolution with formation of the protective BSG layer, and its strong dependence on the LaB6 content during exposure at 1300 °C. One may prefer to add LaB6 rather than La2O3 to the investigated ZrB2-SiC composites, considering that additional amount of B2O3 formed may improve the self-healing ability of the BSG layer developing within the oxide scale by increasing its plasticity and fluidity by decreasing the viscosity. At the same time, it is necessary to limit the decrease in viscosity by proper control of B2O3 content, so that inward diffusivity of oxygen is not increased significantly. An understanding of the effect of LaB6 addition on oxidation resistance of the investigated ZrB2-SiC composites may be considered as a stepping stone for further development of this material in potential high temperature applications involving components in aero-engines and space vehicles.  4.4.  Impact of  the present study  The present study has led to a number of interesting inferences regarding oxidation kinetics and evolution of oxide scale in the investigated ZrB2-SiC based UHTCs having varying LaB6 content, besides generating useful data for isothermal exposure at 1300 °C. Analysis of the results has shown the existence of a direct correlation between the oxidation kinetics and the growth rate of the BSG layer, which primarily comprises the protective part of the oxide scale. As the diffusion of oxygen through the BSG layer is the slowest process in the sequence of events involved in oxidation of the investigated composites, changes in its thickness and viscosity have directly influenced the kinetics of their mass gain. Interestingly, linear relations with slopes having opposite signs have been found for suitably expressing the dependence of the BSG growth rate with LaB6 content for durations of 1−8 h and 8−24 h, Fig. 11. Moreover, for a given composition of the LaB6 containing composite, it takes some time of exposure at high temperature before the protective character is fully established, which coincides with a critical thickness, as well an optimum BSG layer composition and viscosity being reached. As the protective BSG layer appears to be developing continuously in the oxide scale after about an hour of exposure, the oxidation kinetics is slowed until a steady state is reached after duration of 10−12 h. The present study has shown that on exposure at 1300 °C, the least mass gain is observed in the ZSBCL-10 composite after exposure for 8 h, and in the ZSBCL-14 after 24 h, which suggests that the LaB6 content has a significant effect on the kinetics of oxidation and protective scale formation. It has been further inferred that the major cause for change in oxidation behavior with the LaB6 is of course the variation in oxide scale chemistry and resultant viscosity due to varying amounts of La2O3  5009  5. Conclusions  The isothermal oxidation behavior of ZrB2-SiC-LaB6 composites prepared by SPS at 1600 °C has been examined at 1300 °C for 1 h, 8 h or 24 h. Based on the results obtained in the present investigation, the following conclusions can be drawn:  1) The oxidation tests carried out at 1300 °C for 1 h duration have shown random fluctuations in mass change, because a stable and protective BSG layer is not able to develop during the early stage of oxidation. 2) The formation of an appreciable amount of BSG along with a continuous layer of La2Si2O7 on the outermost surface hinders the diffusion of oxygen through the oxide scale; and therefore the mass gain follows approximately parabolic rate law during exposure up to 8 h; whereas it deviates from parabolic to higher-order kinetics for a longer duration of exposure at 1300 °C. 3) The parabolic rate constant is found to scale with the BSG layer thickness, as the diffusion distance for oxygen to reach the oxide scale-composite interface is increased. Decrease in the rate of mass gain with time right after the first hour is suggestive of the growth of the protective BSG layer during this time. 4) The oxidation of ZSBCL-14 has exhibited faster kinetics than the other two investigated composites up to 8 h of exposure, whereas after this period, a significant drop in the rate of oxidation is observed due to the formation of a compact BSG scale with optimum thickness and viscosity, as well as a continuous La2Si2O7 layer on the outermost surface, both of which may have jointly hindered further ingress of oxygen. 5) Microstructural analysis  cross-sections  scale  has  of  the  oxide  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society 40 (2020) 4997-5011  revealed that with an increase in the amount of LaB6, the total thickness of BSG layer reached at the end of 8 h or 24 h is increased, whereas that of the porous ZrO2+BSG layer is reduced. However, the growth rate of the BSG layer during the 8−24 h span of exposure is reduced with respect to that during the 1−8 h time-period, which is suggestive of its protective nature evolving in this period. Both during the initial 1−8 h and later 8−24 h exposures, the growth rate of BSG has shown linear correlation with the amount of LaB6 in the composite, with the slopes of the best-fit lines in these two regimes exhibiting opposite signs.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgements  Technical assistance received from Mr. Mithun Das and Mr. B. Santu Mudliyar, Staff members of Central Research Facility, IIT Kharagpur for characterization of specimens, is gratefully acknowledged.  References  [3]  [6]  [1]  S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. 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},{
  "_id": 116,
  "PDF": "Kinetics and High-Temperature Oxidation Mechanism of Ceramic Materials in the ZrB2–SiC–MoSi2 System.pdf",
  "Text": "['ISSN 1067-8212, Russian Journal of Non-Ferrous Metals, 2018, Vol. 59, No. 1, pp. 76-81. © Allerton Press, Inc., 2018. Original Russian Text © I.V. Iatsyuk, A.Yu. Potanin, S.I. Rupasov, E.A. Levashov, 2017, published in Izvestiya Vysshikh Uchebnykh Zavedenii, Tsvetnaya Metallurgiya, 2017, No. 6, pp. 63-69.  PHYSICAL METALLURGY AND HEAT TREATMENT  Kinetics and High-Temperature Oxidation Mechanism of Ceramic Materials in the ZrB2-SiC-MoSi2 System  I. V. Iatsyuk*, A . Yu. Potanin**, S. I. Rupasov***, and E. A . Levashov****  National University of Science and Technology “MISiS”, Moscow, 119049 Russia *e-mail: ivansvoy@mail.ru **e-mail: a.potanin@inbox.ru ***e-mail: rupasov@misis.ru; vosapur@mail.ru ****e-mail: levashov@shs.misis.ru  Received October 12, 2017; in f inal form, October 20, 2017; accepted for publication October 27, 2017  Abstract ⎯  This study is devoted to the fabrication of the ZrB2-SiC-(MoSi2) compact ceramics according to hybrid technology (self-propagating high-temperature synthesis (SHS) + hot pressing), as well as to investigating its phase composition, structure, and high-temperature oxidation kinetics. Reaction mixtures are prepared according to the following scheme: mechanical activation (MA) of Si + C powders; wet admixing of Zr, B, and Si + C MA-mixture powders; and drying mixtures in a drying oven. The ZrB2-SiC SHS composite powder is formed in a reactor in a combustion mode by elemental synthesis. Compact samples with a homogeneous structure and low residual porosity not exceeding 1.3% are formed by hot pressing the SHS powder. Two compositions are selected for testing, notably, the f irst one calculated for the formation of ZrB2 + 25% SiC; the second composition is similar to the f irst one, but with the addition of 5% of the MoSi2 commercial powder. The microstructure of the samples is presented by dispersed dark gray rounded SiC grains distributed among light faceted ZrB2 grains. The sample with the MoSi2 additive has a more f inely dispersed structure. The high-temperature oxidation of the samples at 1200°C results in the formation of SiO2‒ZrB2-(B2O3) complex oxide f ilms on their surface with a thickness on the order of 20-30 μm, which serve as an eff icient diffusion barrier and lower the oxidation rate. Their structure also contains ZrSiO4 complex oxide after prolonged holding (longer than 10 h). In addition, an insignif icant weight loss of the samples is observed after 10 h  testing, which  is caused by  the volatilization of gaseous oxidation products (B2O2, CO/CO2, MoO3). The sample with the MoSi2 additive shows better resistance to oxidation.  Keywords: zirconium diboride, silicon carbide, hot pressing, kinetics, oxidation, structure  DOI: 10.3103/S1067821218010157  INTRODUCTION  Ceramic ZrB2-based high-temperature materials are characteristic of high melting points (tm > 3245°C), heat and electrical conductivity, and chemical inertness with respect to melts and high thermal-shock resistance [1]. Such a set of unique properties is not inherent  to  traditional high-temperature materials based on Al2O3, ZrO2, and Si3N4. Ceramics based on ZrB2 are considered a promising material for the fabrication of high-temperatures wares, for example, elements of furnaces, plasma-arc electrodes, parts and units of propulsions, and protective coatings operating at temperatures up to 1800°C [1-5]. The main protective mechanism of ZrB2 against oxidation at temperatures above 1000°C is the formation of the ZrO2 barrier layer (tm ~ 2700°C) [6], which has good heat resistance, low vapor pressure, and high mechanical strength. One main disadvantage of the ZrO2 layer is low crack resistance. Alloying ZrB2 with  silicon carbide considerably  increases mechanical strength at high temperatures, as well as wear resistance and heat resistance, and lowers the thermal expansion coeff icient [6-11]. The oxidation product of SiC is silica, which retards further oxidation up to 1700°C. The authors of [12] established that, when adding ~25 wt % SiC into the ZrB2 + SiC composition, the strength characteristics of ceramics formed by hot pressing are minimal and almost do not lower up to the application temperature of 1200°C. The presence of the SiO2 protective f ilm on the carbide surface leads to the interaction following the reaction 2SiO2 + SiC → According to the phase diagram of the Si-O system, monoxide SiO occurs in the solid phase in range t = 1180-1880°C. It follows from this fact that the gaseous phase above SiC will consist preferentially of CO. When oxidizing ZrB2-SiC materials, CO is also able   3SiO + CO.  76  \\x0c', 'KINETICS AND HIGH-TEMPERATURE OXIDATION MECHANISM  77  to worsen the activity of B2O3 and, correspondingly, lower its oxidation probability to B2O2. Oxide B2O3, having low vapor pressures in a temperature range of 1100-1300°C [13, 14], is able to oxidize to volatile B2O2 with the oxygen access [15, 16]. Herewith, the layered structure of the oxidized region is formed, while the formed dense borosilicate scale with the SiO2-ZrO2-B2O3 composition is concentrated preferentially on the material surface [17-20].  Additional alloying with disilicides MoSi2 and ZrSi2 in small amounts (up to 5%) increases the heat resistance of the ZrB2-SiC composition under prolonged holding (higher than 1400°C for 30-50 h) due to  the additional  silicon  source and promotes a decrease in porosity of a compact material after hot pressing [21]. The presence of the disilicide intercrystallite liquid phase during sintering promotes grain ordering, the removal of oxide f ilms, and an increase in the packing density of particles. Disilicide additives also are additional sources of silicon to form SiO2. Metal oxides forming in the structure of a borosilicate scale during the oxidation of silicides also promote the appearance of nonmixing phases with a higher viscosity and low permeability of oxygen, decreasing the intensity of its penetration into the material. MoSi2 itself [22] has high physical properties and heat resistance. However, we should take into account that the disilicide excess can lead to lowering the mechanical properties of ceramics in view of the formation of a large amount of vitreous low-plastic SiO2-B2O3 borosilicate phase at high temperatures, including the sintering process.  One promising method of  fabricating ceramics based on ZrB2 and SiC is self-propagating high-temperature synthesis (SHS) from elements [7, 23, 24], while the main consolidation methods are hot pressing (HP) and spark plasma sintering (SPS). Materials based on zirconium diboride are also in demand as cathode targets for the magnetron sputtering of hightemperature tribological coatings [25].  The authors of [24] considered the peculiarities of combustion and stage character of the structure formation in the combustion wave of the Zr-B-Si-C SHS system. This study is devoted to investigating the kinetics and mechanism of high-temperature oxidation of the ZrB2-SiC-MoSi2 ceramics formed by the hot pressing of SHS powders.  EXPERIMENTAL  We used the following powders as initial reagents:  (i) zirconium of the PTsrK-1 grade with dispersity d = 20 μm;  (ii) silicon formed by the milling of KEF-4.5 single crystals (orientation (100)), d < 45 μm;  d ≤  d ≤  (iii) technical carbon black of the P80 4-T grade,  1 μm; (iv) black amorphous boron of the B-99A grade,  0.2 μm; (v) molybdenum disilicide MoSi2 Technical Specif ications 6-09-03-395-74), d ~ 15 μm. To perform the SHS process, the reaction mixture of powders of Zr, Si, B, and C was prepared, calculated for the formation of 75% ZrB2 + 25% SiC mixture (composition 1) according to the following scheme: (i) mechanical activation (MA) of Si + C powders in a MPP-1 centrifugal planetary mill (CPM) for 2 min with the ratio of weights of the charge and balls of 1 : 15; (ii) wet admixing (in isopropanol) of Zr, B, and Si + C MA mixture powders in a ball mill (BM) for 8 h with the ratio of weights of the charge and balls of 1 : 8. Drying was performed in vacuum at 50°C. Molybdenum disilicide was introduced in an amount of 5% into the ready ZrB2-SiC SHS powder in the BM (composition 2). The synthesis was performed in a SHS-30 reactor in Ar. The product in the form of a porous cake was milled in a BM, after which a powder fraction f iner than 45 μm was isolated. Ceramics was consolidated by hot pressing using a DSP-515 SA press (Dr. Fritsch, Germany) in vacuum at 1800°C, a heating rate of 10 K/min, a pressure of 30 MPa, and isobaric holding for 10 min. Samples in the form of plates 10 × 10 × 5 mm in size were fabricated from compact ceramics for subsequent measurements. Their  faces were preliminarily polished using a Rotopol-21 installation (Struers, Denmark) and subjected to ultrasonic cleaning in isopropanol. The hydrostatic density was determined using a GR-202 AND1 analytical balance (Japan) accurate to 10-4 g, and real density was determined using an AccuPyc 1340 helium pyknometer (Micromeritics, United States). Experiments for high-temperature oxidation were performed in a SShOL 1.1.6/12-M3 electrical furnace in air at 1200°C. The degree of oxidation for 30 h was calculated by weight increment every 0.25, 0.5, 0.75, 1, 2, 3, and 4 h, and then every 5 h. The X-ray phase analysis (XRD) of the initial and oxidized samples was performed using monochro radiation in angle range 2θ  = 10°-110°, matic CuKα and microstructural  investigations were performed using an S-3400N Hitachi scanning electron microscope (Japan) equipped with a NORAN X-ray energy dispersion spectrometer.  RESULTS AND DISCUSSION  The phase composition and residual porosity (Pres) of hot-pressed ceramic samples is presented in Table 1. Main phases are ZrB2 and  low-temperature cubic  RUSSIAN JOURNAL OF NON-FERROUS METALS    Vol. 59    No. 1    2018  \\x0c', '78  IATSYUK et al.  Table 1. Phase composition and porosity of initial ceramic samples  Composition  Pres,  %  ZrB2 (hP3/4)  -SiC (cF8/3)  Si (cF8/1)  MoSi2 (tI6/2)  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  1  2  1.3  1.3  70  68  a = 3.167 c = 3.528  a = 3.155 c = 3.506  28  a = 4.355  2  a = 5.416  -  27  a = 4.352  -  5  a = 3.203 c = 7.841  modif ication β -SiC. In addition, the sample of composition 1 contains a small amount of free silicon, which was also noted in [24] and explained by the partial burning-out of carbon due to high temperatures in the combustion wave. The content of the MoSi2 phase in the sample of composition 2 corresponds to the charged one. The residual porosity of the samples is identical, being 1.3%.  Figure 1 shows microstructures of compact samples. Main structural components are ZrB2 and SiC grains and their agglomerates, which correlates well with the data of semiquantitative XRD. Dispersed dark gray rounded SiC grains are distributed among light faceted ZrB2 grains. The structure of the samples is dense and homogeneous. The sample of composition 2 has a more f inely dispersed structure because the MoSi2 additive during HP forms an intercrystallite phase along the boundaries of ZrB2 grains, which promotes ordering and increasing the packing density of grains [26]. Dependences of varying the sample weight (∆m) during oxidation at 1200°C are presented in Fig. 2. Its most  intense  increment was observed for the f irst hour.  (2)  (3)  (1)  Herewith, SiO2-ZrO2-(B2O3) oxide  f ilms are formed on the surface according to known reactions: ZrB2(s) + 5/2O2(g) →  B2O3(l) + ZrO2(s), SiC(s) + 3/2O2(g) →  SiO2(s) + CO(g), 2SiO2(s) + SiC(s) →  3SiO(s) + CO(g). The formation of the zirconium silicate f ilm is possible for prolonged holding (tens of hours) according to the reaction [27] SiO2(s) + ZrO2(s) →  ZrSiO4(s). The ZrSiO4 compound provides the effect of selfhealing of pores and microcracks in the SiO2-ZrO2- B2O3 borosilicate scale and prevents oxygen diffusion into  the material depth  [27]. However, a  reverse decomposition reaction of ZrSiO4 into SiO2 and ZrO2 is possible in temperature range t = 1200-1500°C [27, 28]. Self-healing also occurs after the melting of B2O3(l) (tm = 450°C) due to the borosilicate formation. An  insignif icant weight  loss of  the samples  is observed after 10-h testing, which is caused by the volatilization of oxidation products B2O2, CO/CO2, and MoO3. Gaseous products can violate the integrity of the oxidized layer, leading to the formation of craters.  (4)  (а)  20 μm  (b)  20 μm  Fig. 1. Microstructure of compact ceramics of compositions (a) 1 and (b) 2.  RUSSIAN JOURNAL OF NON-FERROUS METALS    Vol. 59    No. 1    2018  β \\x0c', 'RUSSIAN JOURNAL OF NON-FERROUS METALS    Vol. 59    No. 1    2018  KINETICS AND HIGH-TEMPERATURE OXIDATION MECHANISM  79  An increase in temperature to 1100°C decreases the protective properties of the B2O3 melt layer because of its oxidation to gaseous B2O2, which partially evaporates [13-16], thereby violating the integrity of the SiO2-ZrO2-B2O3 protective f ilm. Molybdenum disilicide starts to oxidize at  t = 1200°C with the formation of volatile MoO3 oxide and Mo5Si3 and SiO2 solid f ilms according to the following reactions [22]: 2MoSi2(s) + 7O2(g) → 5MoSi2(s) + 7O2(g) → According to the data of semiquantitative XRD, the oxidized layer after 30 min contains ZrO2, β -SiC,   2MoO3(g) + 4SiO2(s),  (5)   Mo5Si3(s) + 7SiO2(s).  (6)  and ZrB2 (Table 2). A low content of MoSi2 is f ixed for the sample of composition 2. In addition to listed phases, the oxidized layer after 30-h testing also contains oxides ZrSiO4 and SiO2.  Figure 3 shows microstructures of fractures near the surface of the samples oxidized at t = 1200°C for 30 h. A layer with a thickness of 10-15 μm consisting of ZrO2 grains of about 5 μm in size is formed on the surface of the sample of composition 1 (Fig. 3a) after testing in mentioned conditions. The space between the grains and ZrO2 agglomerates is f illed with SiO2 bound into a borosilicate scale. A thin porous sublayer consisting of ZrO2 grains 1-5 μm in size and coarser grains and ZrSiO4 agglomerates is arranged under this  Fig. 2. Variation in the weight of the samples of compositions 1 and 2 during the oxidation at 1200°C for (a) 30 h and (b) the f irst hour of testing.  Δm/S, mg/mm2 0.0 40 0.035 0.030 0.025 0.020 0.015 0.010 0.005 0  0.005  0.010  0.015  0.020  0.025  Δm/S, mg/mm2 0.030  (а)  35 τ, h  5  10  15  20  20  0  0  (b)  1.0 τ, h  0.1  0.2  0.3  0.4  0.5  0.6  0.7  0.8  0.9  0  1  1  2  2  Table 2. XRD results of the sample surface after testing for heat resistance for 30 min and 30 h at t = 1200°C  Phase (Pearson  symbol)  Composition 1  Composition 2   = 30 min  30 h  30 min  30 h  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  wt %  lattice  parameter, Å  ZrO2 (mP12/3)  74  a = 5.142 b = 5.196 c = 5.317  = 99.181  12  a = 5.139 b = 5.208 c = 5.307  = 98.850  52  a = 5.143 b = 5.194 c = 5.318  = 99.206  13  a = 5.152 b = 5.192 c = 5.320  = 99.059  ZrB2 (hP3/4)  12  a = 3.153 c = 3.50 4  2  -  25  a = 3.166 c = 3.526  7  a = 3.168 c = 3.527  -SiC (cF8/3)  14  a = 4.352  8  a = 4.358  21  a = 4.351  10  a = 4.352  SiO2 (tP12/1)  -  30  a = 4.988 c = 6.923  -  18  a = 5.011 c = 6.945  ZrSiO4 (tI24/3)  -  47  a = 6.605 c = 5.976  -  52  a = 6.605 c = 5.976  MoSi2 (tI6/2)  -  -  2  -  -  τ β β β β β \\x0c', '80  IATSYUK et al.  (а)  20 μm  (b)  20 μm  Fig. 3. Microstructures of fractures of the samples of compositions (a) 1 and (b) 2 in the surface layer region oxidized at t = 1200°C for 30 h.  layer. The thickness of this sublayer is no larger than 10 μm. Pores formed during the material oxidation can be f illed with gaseous products of oxidation of SiC (SiO and CO) diffusing from the reaction zone to the surface. The ZrB2-SiC layer depleted with respect to SiC is arranged further, and unoxidized initial material is arranged below. The total thickness of the oxidized layer is 20-30 μm. The SiO2 layer with inclusions of small ZrO2 grains no more than 5 μm in size is retained on the surface of the sample with the MoSi2 additive after 30-h testing (Fig. 3b). An oxidized layer is rather dense; no cylindrical or through pores are observed. Its thickness compared with the sample without MoSi2 additive increases approximately twofold, being 20-25 μm. It is very diff icult to determine the locality of the phase of the ZrSiO4 ternary oxide in view of a broad excitation region of the electron microscope detector. An unoxidized initial material is arranged under this layer. Thus, ceramics with the addition of MoSi2 showed the best oxidation resistance, which is caused by the larger fraction of silicon-containing phases. It is seen from Fig. 2  that  the  smaller weight  increment  is observed for the f irst 10 h of oxidation and an insignif icant weight decrement is observed after 10 h of testing.  CONCLUSIONS  The oxidation kinetics and mechanism of hotpressed ceramics of two compositions ZrB2-SiC and ZrB2-SiC-MoSi2 are investigated in our study. It is shown that the protective f ilm of zirconium silicate ZrSiO4 and borosilicate scale SiO2-ZrO2-B2O3 20- 30 μm in thickness, preventing the oxygen penetration into the material depth, are formed on the surface of the  samples. 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Parthasarathy, T.A., Rapp, R.A., Opeka, M., and Cinibulk, M.K., Modeling oxidation kinetics of SiC-containing refractory diborides, J. Am. Ceram. Soc., 2012, vol. 95, no. 1, pp. 338-349. 19. Eakins, E., Jayaseelan, D.D., and Lee, W.E., Toward oxidation-resistant ZrB2-SiC ultra high temperature  ceramics, Metal. Mater. Trans. A, 2011, vol. 42, no. 4, pp. 878-887. 20. Gao, D., Zhang, Y., Fu, J., Xu, C., Song, Y., and Shi, X., Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: Formation of zircon, Corros. Sci., 2010, vol. 52, no. 10, pp. 3297-3303. 21. Silvestroni, L. and Sciti, D., Effects of MoSi2 additions on the properties of Hfand Zr-B2 composites produced by pressureless sintering, Scr. Mater., 2007, vol. 57, pp. 165-168. 22. Potanin, A.Yu., Pogozhev, Yu.S., Levashov, E.A., Novikov, A.V., Shvindina, N.V., and Sviridova, T.A., Kinetics and oxidation mechanism of MoSi2-MoB ceramics  in  the  600-1200°C  temperature  range, Ceram. Int., 2017, vol. 43, no. 13, pp. 10 478-10 486. Iatsyuk,  I.V.,  Pogozhev, Yu.S., Levashov, E.A., Novikov, A.V., Kochetov, N.A., and Kovalev, Yu.D., Peculiarities of production and high-temperature oxidation of SHS ceramics based on zirconium boride and  23.  24.  silicide, Izv. Vyssh. Uchebn. Zaved., Poroshk. Metal.  Funkts. Pokr yt., 2017, no. 1, pp. 29-41. Iatsyuk, I.V., Pogozhev, Yu.S., and Novikov, A.V., Synthesis of ZrB2-SiC high-temperature ceramics in the combustion mode, Tsvet. Met., 2017, no. 12, pp. 71-77. 25. Kiryukhantsev-Korneev, P.V., Lemesheva, M., Yatsyuk, I., Shtansky, D.V., and Levashov, E.A., Comparative investigation of Zr-B-(N), Zr-Si-B-(N), and Zr-Al-Si-B-(N) hard coatings, 44th Int. Conf. on  Metallurgical Coatings and Thin Films   (ICMCTF),  San Diego, California: 2017, pp. 50-51. 26. Sciti, D., Guicciardi, S., and Bellosi, A., Properties of pressureless-sintered ZrB2-MoSi2 ceramic composite, J. Am. Ceram. Soc., 2006, vol. 7, pp. 2320-2322. 27. Yu, Y., Luo, R., Xiang, Q., Zhang, Y., and Wanga, T., Antioxidation properties of a BN/SiC/Si3N4-ZrO2- SiO2 multilayer coating for carbon/carbon composites, Surf. Coat. Technol., 2015, vol. 277, pp. 7-14. 28. Liu, J., Cao, L.-Y., Huang, J.-F., Xin, Y., Yang, W.-D., Fei, J., and Yao, C.-Y., A ZrSiO4/SiC oxidation protective coating for carbon/carbon composites, Surf. Coat. Technol., 2012, vol. 206, pp. 3270-3274.  Translated by N. Korovin  RUSSIAN JOURNAL OF NON-FERROUS METALS    Vol. 59    No. 1    2018  α \\x0c']"
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  "_id": 117,
  "PDF": "Kinetics-and-evolution-of-oxide-scale-during-various-stages-of-isothermal-oxidation-at-1300C-in-spark-plasma-sintered-ZrB--SiC--LaB-composites2020Journal-of-the-European-Ceramic-Society.pdf",
  "Text": "['Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e lsev ie r .com / loca te / jeu rce ramso c  Original Article  Kinetics and evolution of oxide scale during various stages of isothermal oxidation at 1300 °C in spark plasma sintered ZrB2 - SiC - LaB6 composites  Sunil Kumar Kashyapa,*, Ankit Kumarb, Rahul Mitraa  a Department of Metallurgical and Materials Engineering, b Department of Metallurgical and Materials Engineering,  Indian Institute of Technology Kharagpur, Kharagpur 721302, West Bengal, India Indian Institute of Technology Roorkee, Roorkee 247667, Uttarakhand, India  A R T I C L E  I N F O  Keywords: ZrB2-SiC-LaB6 composites Ultra-high temperature ceramics Spark plasma sintering Isothermal oxidation Kinetics  A B S T R A C T  The kinetics and oxide scale evolution during isothermal exposure of spark-plasma sintered ZrB2-20 SiC-LaB6 (7, 10 or 14 vol%) composites at 1300 °C for 1, 8, or 24 h have been examined. Random mass change observed during the first hour stems from a non-protective scale. The variation of mass gain with time is expressed by near-parabolic rate law during 0−8 h time period, and by relations indicating slower kinetics during 8−24 h, with parabolic rate constants (kp) decreasing sharply. Microstructural examination has shown a continuously evolving layered oxide scale comprising La2Si2O7, borosilicate glass (BSG), ZrSiO4, and ZrO2, where the BSG layer growth rate scales linearly with LaB6 content during 0−8 h, but shows an opposite trend on further exposure. During both time periods (0−8 and 8−24 h), kp decreased following a linear relationship with increasing BSG layer thickness, indicating its key role as diffusion barrier for oxygen.  1.  Introduction  Ultra-high temperature ceramics (UHTCs) developed in the 1960s, are known for their application in thermal insulation tiles or thermal barrier coatings, because of their high melting points, high thermal conductivity, high ablation resistance, and chemical attack resistance [1,2]. Zirconium diboride is one of the materials, which has been extensively considered for use as a UHTC [3]. Among the various UHTCs including borides, nitrides, and carbides, the use of ZrB2 is preferred due to its moderate density (6.1 g/cc) and high melting temperature (3250 °C), which favors the application of these materials in leadingedges of re-entry vehicles and rocket nozzles [4,5]. As reported in the literature, it is very difficult to densify monolithic zirconium diboride due to strong covalent bonding, and therefore to deal with such a problem, various types of reinforcements or additives have been used to promote densification and also lower the process temperature [6,7]. Generally, carbides (SiC, B4C), nitrides (AlN, Si3N4), and di-silicides (MoSi2, TaSi2) have been used as sintering aids for improving the densification of ZrB2 [8-13]. For the UHTCs, to sustain their structural integrity during the functioning, high oxidation resistance is considered as a pivotal requirement. Earlier studies have confirmed that the ZrB2 -SiC composites possess high oxidation resistance due to the formation of a protective outer layer of borosilicate glass (BSG) in the temperature range of 1000 °C-1700 °C. With the formation of an outer BSG layer, an  inner SiC-depleted region below the intermediate ZrO2-SiO2 layer is formed inside the oxide scale of the composite, as expected on considering the thermodynamic stability of the oxidation products [14-16]. The improved oxidation resistance is facilitated by the formation of an adherant and compact oxide scale enriched in BSG, which acts as an effective barrier against the inward diffusion of oxygen [17,18]. Along with improving the oxidation resistance, the presence of SiC particles also resists the grain growth of the ZrB2 phase during sintering, with benefits of superior thermo-mechanical properties as compared to the boride alone [19]. Some of the earlier works on the ZrB2 based composites have shown that the isothermal oxidation tests carried out for reasonably long durations exhibit mass gain kinetics following a parabolic rate law at temperatures < 1100 °C; whereas deviation towards a near-linear relationship has been observed for oxidation kinetics at higher temperatures [15,20-22]. However, the addition of SiC is found to change the mass gain kinetics from liner to parabolic in the temperature range of 1300−1600 °C by providing an effective barrier for oxygen diffusion through the oxide scale [14,23]. Liu et al. [24] have reported that the oxidation kinetics of ZrB2-20 vol% SiC composite follow a parabolic rate law; whereas on addition of 20 vol% ZrC at the expense of ZrB2, it shows near-linear kinetics at 1600 °C. Furthermore, the ZrB2-30 vol% SiC composite has followed the cubic rate law at 1500 °C due to increased viscosity along with reduction in the amount of ZrO2 inclusions  ⁎ Corresponding author. E-mail address: sunil.kashyap90@gmail.com (S.K. Kashyap).  https://doi.org/10.1016/j.jeurceramsoc.2020.07.053 Received 7 June 2020; Received in revised form 17 July 2020; Accepted 20 July 2020 0955-2219/ © 2020 Elsevier Ltd. All rights reserved.  Please cite this article as: Sunil Kumar Kashyap, Ankit Kumar and Rahul Mitra, Journal of the European Ceramic Society,  https://doi.org/10.1016/j.jeurceramsoc.2020.07.053  \\x0c', \"(4.0 ± 1.4 μm), LaB6 (7.9 ± 2.8 μm), and B4C (1.84 ± 1.3 μm) were used to prepare the composites for this study. The purity of all the powders was ≥99.5 %. The powders of ZrB2 and SiC were obtained from the H.C. Starck, GmbH, Goslar, Germany, whereas the LaB6 and B4C powders were purchased from the Alfa Aesar, Massachusetts, USA, and Boron Carbide India Limited, Mumbai, India, respectively. Furthermore, phenolic resin (phenol-formaldehyde with ethanol as solvent), type ABRON PR100 obtained from ABR Organics Limited, Hyderabad, India, has been used as a source of carbon in the raw materials. The carbon content of the phenolic resin has been found to be approximately 37 wt% by thermogravimetric analysis. Three different composite powder mixtures (ZSBCL-7, ZSBCL-10, ZSBCL-14) having compositions of ZrB2 + 20 vol % SiC + 5.6 vol % B4C + 4.8 vol % Carbon, with varying vol% of LaB6 (7%, 10 % and 14 %) were prepared. These powders were mixed using a planetary mono-mill (model Pulverisette 6, Fritsch GmbH, Idar-Oberstein, Germany) with acetone as the medium using vials and balls of WC-Co. After the completion of milling for 6 h at 250 rpm, the acetone was drained out, and the blended powders were dried at 350 °C for 2 h in the furnace. The composite powders were consolidated into pellets of 20 mm diameter by spark plasma sintering (Dr. Sinter, Japan) for 7 min at 1600 °C with application of ram pressure equal to 50 MPa. The bulk density of the sintered composites was measured by Archimedes' principle, whereas the theoretical densities were calculated by the rule of mixture method. Samples with dimensions of 6 mm x 4 mm x 4 mm were cut from the pellets of spark plasma sintered (SPS) composite using EDM wire-cutting, and then polished on emery paper, followed by cloth polishing with 6 μm diamond paste. Each sample was ultrasonically cleaned in acetone, and the constituent phases having different compositions were identified with the help of X-ray diffraction using a Cu target (wavelength of Cu Kα radiation = 0.15418 nm) on a high resolution diffractometer (Bruker D8 Discover, Germany) operated at 0.02˚/s step size. Microstructures of all the composites were examined by a field emission scanning electron microscope (Merlin FESEM, Carl Zeiss, Oberkochen, Germany).  2.2. Oxidation tests  Isothermal oxidation tests were performed on the composites with aforementioned compositions for 1 h, 8 h, and 24 h at 1300 °C using an electric vertical chamber furnace (designed and manufactured by BYSAKH & COMPANY, Kolkata, India) with arrangement of resistive heating by SiC elements. The change in mass during isothermal exposure was measured at selected intervals by a mechanical balance  in BSG, as reported in an earlier study [23]. The effect of the addition of La2O3 on ZrB2-20 vol% SiC composites has been investigated by Zapata-Solvas in the temperature range of 1400−1600 °C [25]. In this study, it has been found that in this temperature range, the ZrB2-20 vol% SiC composites exhibit kinetics of mass gain following a power law (Δwn = kt), where 1 ≤ n ≤ 2, indicating oxidation rates to be intermediate between those expected from linear and parabolic kinetics; whereas the addition of La2O3 to the composite has led to the kinetics based on the parabolic rate law, indicating a significant increase in oxidation resistance within the same temperature range [25]. Interestingly at 1600 °C, the oxidation exponent has been found to be 8, which is attributed to the formation of new protective (Zr/Si)OxCy phase. The addition of LaB6 to the ZrB2-SiC composites is known to contribute to the oxidation resistance by forming a refractory La2Zr2O7 and/or La2Si2O7 scale on the outermost surface of oxide scale [6,17,26,27]. Moreover, the increase in B-content due to LaB6 addition is expected to contribute an additional amount of B2O3, which in turn would enhance the fluidity and consequently, the self-healing capacity of the BSG-rich scale. Zhang et al. [26] have performed ablation tests at 2400 °C by using an oxyacetylene torch and reported that the addition of 10 vol% of LaB6 to the ZrB2-SiC composites remarkably enhanced the oxidation resistance by forming a compact La2Zr2O7 scale. Further, the oxidation resistance of the same composition was investigated by Jayaseelan et al. [27] at 1600 °C for 1 h in air. They have reported that the oxygen diffusivity through the La2Zr2O7 scale is lower than that of ZrO2, but it is higher as compared to that in the BSG. Both high viscosity and high melting temperature of the La2Zr2O7 oxide formed on the LaB6 containing composites during high temperature exposure, can be considered as the source of motivation to consider these composites as the suitable candidate material for the hypersonic applications. On the other hand, Monteverde et al. [17] have reported the absence of La2Zr2O7 on the outermost surface of the oxide scale during the arc-jet test of ZrB2-SiC-LaB6 composites at 1700 °C. The lack of La2Zr2O7 has been attributed to the high heat flux and a short duration of exposure, which suppresses the zirconate formation, thereby reducing the oxidation resistance. Further, it has been reported that the addition of LaB6 is detrimental, since it reduces the eutectic temperature of the oxide scale, and increases the oxygen vacancy concentration [6]. Although the oxidation behavior of ZrB2-SiC composites has received quite extensive attention, information about the effect of rare earth compound addition on the kinetics of oxidation is scarce [25]. Moreover, most of the earlier studies on the addition of LaB6 to the ZrB2-SiC composites have been investigated for exposure temperatures ≥1600 °C, none of those have focused on the kinetics of oxidation. Moreover, very limited understanding is available about early and intermediate stages of oxidation, which are known to control the later stages of oxide scale formation. Furthermore, the mass change kinetics and its relationship with the oxide scale evolution during initial and intermediate durations of exposure at a given temperature also needs to be examined for understanding how the mechanisms of protection against oxidation evolve with time. In view of the aforementioned aspects, this paper reports the results of investigation on the effects of varied LaB6 additions at various stages of oxidation including early, intermediate, and after exposure for 24 h, with emphasis on the evolution of the oxide scale at each stage. For this purpose, the ZrB2-20 vol % SiC composites processed by spark plasma sintering with B4C and C as additives, and having 7, 10 or 14 vol% LaB6 have been isothermally exposed for 1 h, 8 h, or 24 h at 1300 °C in dry air, and the evolution of oxide scale has been investigated.  2. Experimental Procedure  2.1. Preparation and characterization of composite  Commercially  available  powders  of  ZrB2  (5.4 ± 2.2 μm),  SiC  \\x0c\", 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 1. SEM images showing the microstructure of spark plasma sintered (a) ZSBCL-7, (b) ZSBCL-10, and (c) ZSBCL-14 composites; as well as (d) the presence of LaBO3 in ZSBCL-14 composite along with (e) EDS point analysis spectra showing the peaks from the constituent elements present in all the phases.  microstructures and compositions was carried out by FESEM and EDS. As a part of this effort, the cross-sections of the oxide scales were examined by FESEM with EDS to identify the composition beneath the outer surface till the oxide-composite interface.  3. Results  3.1. Microstructural observation  The density of ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites was measured by Archimedes principle, and the corresponding values of relative density (ratio of experimentally measured density to the theoretical density calculated on the basis of rule of mixtures) were found to be 98 %, 99 %, and 99.5 %, respectively, as reported in the author’s previous work [28]. Furthermore, the XRD patterns of the composites sintered by SPS at 1600 °C with 50 MPa pressure have shown the presence of ZrB2, SiC, and LaBO3 phases, which were found to be similar to those sintered at 1800 °C with 50 MPa pressure, as discussed in an earlier publication [29]. The SEM images of the metallographically polished surfaces depicting the microstructures of the ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites as shown in Fig. 1(a)-(c), respectively, exhibit a reasonable uniformity in the distribution of ZrB2, SiC, and B4C phases in the composites. Furthermore, the presence of LaBO3 is visible in the magnified image of the ZSBCL-10 composite, as shown in Fig. 1(d). The presence of these phases has been qualitatively confirmed by EDS elemental analysis, as shown in Fig. 1(e). The average ZrB2 grain sizes in ZSBCL-7, ZSBCL-10 and ZSBCL-14 composites processed by SPS at 1600 °C have been found to be 2.6 ± 0.6 μm, 3.7 ± 0.5 μm, and 2.7 ± 0.6 μm, respectively, as reported in a previous study [28]. The amount of SiC in the microstructures of the sintered composites [Fig. 1(a-c)] is found to be more or less unchanged at 20 vol% irrespective of the LaB6 content, as confirmed through quantification by  image analysis with the help of Image-J software. Although it was possible to identify the B4C occasionally in the microstructure by EDS analysis, yet its peaks could not be detected in the XRD patterns from the sintered composites, because its amount was low (< 5 vol%) due to its partial consumption for reduction of the oxide impurities [31]. Moreover, the variation between the distributions of SiC and B4C in the investigated composites with different amounts of LaB6 has been found to be quite insignificant. Small peaks of LaBO3 could be observed in the XRD patterns from the as-sintered ZSBCL-7, ZSBCL-10 and ZSBCL-14 composites, with their intensities scaling with the LaB6 content, as has been reported in an earlier study by the authors [29]. Formation of LaBO3 in the microstructure has been ascribed to a part of LaB6 having reacted with the surface oxides of powders. Such scavenging of oxygen is found to have a significant role in the process of densification during sintering [28,29]. Consumption of LaB6 to form LaBO3 may be considered to have lowered the amount of former compound in the microstructures of the sintered composites. Quantitative image analysis of the SEM (BSE) images depicting the microstructures, carried out by using the Image-J software on the basis of atomic number contrast has shown the amount of LaBO3 to be in the range of 5−7 vol% in case of the investigated composites. Although B2O3 and SiO2 were present as impurities in raw materials, their amounts have been found to be significantly reduced in the sintered composites due to reduction of these oxides by B4C and C, as well as scavenging action of LaB6 [28]. Moreover, in tune with this observation, the ZrB2 matrix grain boundaries as well as ZrB2-SiC interfaces in the investigated composites have been found to be sharp and free of any glassy phase.  3.2. Oxidation behavior  3.2.1. Mass change The variation  3  of mass  change  observed  in  the  investigated  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 2. Plots of change in mass with respect to time at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h isothermal duration.  composites with respect to time, upon isothermal exposure for 1 h, 8 h or 24 h duration at 1300 °C, are shown in Fig. 2(a)-(c), respectively. Isothermal exposure of the Ni-Chrome and platinum wires at the abovementioned temperatures have shown very negligible mass change per unit surface area compared to that observed in case of the composites [18]. The results depicted in Fig. 2(a) show a significant amount of fluctuation in the trend observed for mass change within the first hour of exposure, probably because a protective scale is not able to develop during the early stage of oxidation. Interestingly, the plots in Fig. 2(b) indicate a higher mass gain for the ZSBCL-14 than that recorded for the ZSBCL-10 composite, whereas those in Fig. 2(c) indicate an opposite trend with respect to that observed after exposure for 8 h. The results in Fig. 2(c) show that with increase in the duration of exposure up to 24 h, the net mass gain experienced by the ZSBCL-14 composite is less with greater flattening of the curve, when compared with that of the ZSBCL10. However, it is clear from the graphs in both Fig. 2(b) and (c) that the ZSBCL-7 composite shows a higher mass gain than the other two composites at both the oxidation temperatures for any of the investigated isothermal durations.  3.2.2. Phase identification of oxides scale by XRD analysis The X-ray diffraction (XRD) patterns obtained from the oxide scales formed on the samples of ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites exposed at 1300 °C for 1 h, 8 h, and 24 h duration, are shown in Fig. 3(a)-(c), respectively. The XRD pattern from the oxide scale formed after exposure at 1300 °C for 1 h (Fig. 3(a)) shows the peaks corresponding to ZrO2 (00-037-1484), La2O3 (00-022-0369), and La2Si2O7 (01-072-2456). On the other hand, the XRD patterns from the oxide scale formed after exposure at 1300 °C for 8 h or 24 h show an additional peak of ZrSiO4 (01-083-1374) along with all the aforementioned peaks, but with the exception of La2O3. The difference in between the peak intensities of a given phase in the XRD patterns obtained from the oxide scales formed on the composites exposed at 1300 °C for 8 h or 24 h as shown in Fig. 3(b) and (c), respectively, might be due to the changes in the relative amounts of phases present in the oxide scales, as well as variations in oxide scale thickness and surface roughness altering the depth of location of the crystalline oxide constituents.  3.2.3. Study of oxide scale by SEM 3.2.3.1. Oxide surface morphology. The SEM images depicting the top surfaces of the oxide scales formed on the ZSBCL-7, ZSBCL-10, and ZSBCL-14 composites exposed at 1300 °C for 1 h are shown in Fig. 4(a)-(c), respectively. The magnified views of the selected locations in the aforementioned SEM images have been also incorporated as inset. It is apparent that these surfaces are rough and porous. Comparison of the low magnification SEM images in Fig. 4, depicting the oxide scales formed on the ZSBCL-14 is more compact and less porous compared to that of ZSBCL-7 or ZSBCL-10. The results of EDS analysis, as shown in Fig. 4(d), along with the XRD results from the oxide scales formed after oxidation for 1 h have shown evidence for the formation of ZrO2, ZrSiO4, and La2Si2O7. Additionally, the presence of SiO2 phase has been confirmed by EDS analysis, and this inference is further substantiated by the presence of a hump at diffraction angles (2θ) (≤35° for 8 h exposure and ≤40° for 24 h exposure) in the XRD patterns, indicating the presence of amorphous (glassy) phase. Interestingly, such a hump is not distinct in the XRD patterns from the oxide scales formed on the composites exposed at 1300 °C for 1 h, probably because the amount of amorphous phase is very low. The presence of ZrSiO4 phase in the oxide scale formed on exposure at 1300 °C in the present study is in agreement with the results reported in an earlier study [30]. The FESEM images of the samples exposed at 1300 °C for 8 h or 24 h show the formation of a dense scale comprising La2Si2O7 appearing to be clustered or forming a network in the BSG matrix, as shown in Figs. 5 and 6. In each of these figures, ZSBCL-7, ZSBCL-10, and ZSBCL-14 are represented as (a), (b), and (c), respectively. It can be inferred from these images that the formation of the La2Si2O7 phase becomes more predominant in comparison to other oxide scale constituents, as the LaB6 content of the composite is increased to ≥10 vol%. Whereas the oxide scale formed after exposure at 1300 °C for 8 h on the ZSBCL-7 has exhibited the presence of La2Si2O7 phase as islands (Fig. 5(a)), those on the composites with higher LaB6 content appear to be in the form of large clusters or networks. Furthermore, these small islands have combined during 24 h of isothermal exposure to form big clusters of La2Si2O7, whereas ZrSiO4 has formed due to reaction between SiO2 and ZrO2 at the surrounding locations, as shown in Fig. 6(a). Intuitively, the  Fig. 3. XRD patterns obtained from the oxide scales formed during isothermal oxidation tests in the air at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h duration.  4  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 4. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 1 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites; as well as (d) EDS point analysis spectra showing the peaks from the constituent elements present in all the phases formed after oxidation.  amount of La2Si2O7 phase formed has increased with increase in the duration of isothermal exposure from 1 h to 24 h. The XRD patterns of oxide scale (Fig. 3) have also shown the intensity of La2Si2O7 peaks to scale with both isothermal test duration, and LaB6 content of the composite. The absence of peaks belonging to La2O3 in the XRD patterns from the oxide scales formed by exposure for 8 h or 24 h at 1300 °C is ascribed to their consumption in the reaction with SiO2 to form La2Si2O7.  3.2.3.2. Cross-section of oxide scale. A typical SEM (BSE) image depicting the cross-section of the oxide scale formed on the ZSBCL-14 composite due to the exposure at 1300 °C for 1 h is shown in Fig. 7(a), whereas the EDS X-ray maps depicting the locations of enrichment of Zr, Si, O, La, and B are shown in Fig. 7(b-f), respectively. Examination of the aforementioned EDS X-ray maps as shown in Fig. 7(b-f) indicates the formation of an oxide scale comprising ZrO2-B2O3-SiO2 mixture.  Typical backscattered electron (BSE) micrographs depicting the crosssections of the oxide scale formed by exposure at 1300 °C for 1 h on the ZSBCL-14 at higher magnification indicates that the oxide scale comprises a SiO2-rich outer layer of BSG with inclusions of ZrO2 having a greater area fraction near the interface of this layer with a porous inner layer composed of a ZrO2+SiO2-B2O3 mixture. It is interesting to note that no La-containing oxide is visible as an outer layer in the cross-section image of the oxide scale formed after 1 h duration, and this could be because of its presence as discrete islands with varied sizes, and not as a continuous film at the top surface depicted in Fig. 4(c). On examination of the SEM images in Fig. 8(b) and (c) showing the oxide scales formed on exposure at 1300 °C for 8 h and 24 h, respectively, it is evident that the outer layer is made of a thin but more or less continuous La2Si2O7, followed by inner layers comprising the following oxides sequentially arranged: BSG, ZrO2, and porous ZrO2 with traces of BSG. Further, the thickness of BSG  Fig. 5. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 8 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites.  5  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 6. FESEM images showing the top surfaces of the oxide scales formed after isothermal oxidation at 1300 °C for 24 h in the air on (a) ZSBCL-7, (b) ZSBCL-10, (c) ZSBCL-14 composites.  layer appears to have increased with increasing duration of the isothermal hold at 1300 °C. Based on the comparison of the SEM images depicting the cross-sections of the oxide scales formed at 1300 °C for durations from 1 to 24 h, it is appropriate to infer that the growth of a distinct protective outer layer of La2Si2O7 over the composites is possible only after a minimum duration, which may be related to their LaB6 content. The thicknesses of the oxide scales formed on the investigated composites after exposure for 1 h, 8 h, and 24 h at 1300 °C are presented in Table 1. Based on the results shown in this table, the following inferences may be drawn: (i) on exposure for 1 h or 8 h, the thickness of the oxide scale formed on the ZSBCL-14 is the highest, followed by those on ZSBCL-7 and ZSBCL-10; and (ii) thicknesses of the ZrO2 layer in the oxide scales of the investigated composites are found to be comparable on exposure for 8 h or 24 h; whereas (iii) thickness of BSG layer in the oxide scales increases in the following order ZSBCL7 < ZSBCL10 < ZSBCL14. In other words, after 24 h of exposure, the ZSBCL-14 exhibited the lowest oxide scale thickness with maximum BSG formation amongst the investigated composites. The lowest total oxide scale thickness observed in case of the ZSBCL-10 after 1 or 8 h, and of ZSBCL-14 after 24 h of isothermal exposure appear to be consistent with their lowest mass gain among the investigated composites during isothermal hold for the respective durations as shown in Fig. 2.  4. Discussion  4.1. Thermodynamic parameters of oxidation reactions  The thermodynamic aspects of the oxide scale formation during isothermal exposure at 1300 °C have been analyzed. The values of change in free energy of formation per mole of oxygen [ΔGf (kJ/mol)], the equilibrium constant (Kf,eq), and the partial pressure of oxygen ( ) have been calculated with the help of data available in the literature [31-33]; and presented in Table 2. The values of Kf,eq have been calculated by using the following relation:  ΔGf = -RT ln (Kf,eq)  (1)  where R is the universal temperature. Further, the values of for the formation of oxides have been calculated by substituting the values of Kf,eq in the following equations:  constant,  and T is  the  absolute  ideal  gas  (2)  (3)  Fig. 7. Cross-section of oxide scale formed on ZSBCL14 by exposure for 1 h at 1300 °C: (a) SEM (BSE) image, and EDS X-ray maps showing enrichment of (b) Zr, (c) Si, (d) O, (e) La, and (f) B.  6  pO2pO2=()()()Kaaap**()feqZrOZrOBOZrBO,()2/52/52/5222322=()Kapap*()()*()feqSiOSiOCOSiCO,()2/32/32/3222\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 8. Cross-section of oxide scale formed on ZSBCL14 by exposure at 1300 °C for (a) 1 h, (b) 8 h, and (c) 24 h duration.  (4)  (5)  where is the activity of various reactants and products mentioned in the considered reactions as outlined in Table 2, and is the equilibrium partial pressure of oxygen for the given reactions. The values of have been calculated by taking the activity of pure substance as unity, and ignoring the presence of CO in Eq. (5) [15]. The results in Table 2 indicate that all the reactions are feasible at 1300 °C, because the ΔGf (kJ/mol) values are negative. Based on the ΔGf (kJ/mol) values shown in Table 2, the chemical stability of oxides at investigated temperatures is found to decrease in the following order: La2Si2O7 > ZrSiO4 > ZrO2 > SiO2 > SiO > La2O3. The evidence for the presence of ZrSiO4 in the oxide scales as depicted in Figs. 3, 4, and 6 may be considered to be consistent with its co-existence in thermodynamic equilibrium with ZrO2, as predicted by the ZrO2-SiO2 phase diagram [34]. The absence of La2Zr2O7 may be attributed to the significantly lower (more negative) ΔGf (kJ/mol) value of La2Si2O7 (-2809 KJ/mol) as compared to that of the former oxide (-150 KJ/mol) at 1300 °C [35]. The lower free energy of formation of La2Si2O7 may have preferably promoted its formation, rather than that of La2Zr2O7 during exposure at 1300 °C. It may be noted that a few earlier studies [6,17,26,27] have reported the presence of the La2Zr2O7 in the oxide scales of ZrB2-SiC-LaB6 composite exposed only at temperatures ≥1600 °C, and therefore its absence from the oxide scales formed at 1300 °C in the present study is not surprising. Based on the data presented in Table 2, the free energy of the Reaction (2) is more negative than the Reaction (4), which indicates that the formation of SiO2 is more strongly favored compared to that of SiO at the investigated temperatures. The formation of SiO at temperatures > 1500 °C by active oxidation of SiC causing the formation of a SiC-depleted region, at the oxygen-deficient locations (or regions having lower partial pressure of oxygen) underneath the top surface, has been reported by many researchers [15,36-38]. However, a few authors have reported that the SiC-depleted region is formed even at a  of oxygen  low partial pressure  lower temperature 1200 °C under [39,40]. A high value of for La2O3 formation indicates a relatively lower driving force for the oxidation reactions leading to its formation, which can be easily correlated with limited growth of the La-based oxide scale (as shown in Fig. 8) during oxidation. It has been shown in earlier reports that the oxidation of LaB6 starts at 700−800 °C, resulting in the formation of a few intermediate oxides, and those oxides may convert to La2O3 at temperature ≥1200 °C [21,41]. On the other hand, the oxidation of ZrB2 and SiC starts at ≥800 °C and ≥1100 °C, respectively [15].  4.2. Oxidation kinetics  The oxidation kinetics of the composites have been analyzed on the basis of the data representing the change in mass (ΔW) by oxidation during isothermal exposure (t) to calculate the oxidation rate-related parameters such as the oxidation exponent (n), the general rate constant (k), and the parabolic rate constant (kp). The oxidation kinetics can be determined by using the power-law equation:  (ΔW)n = kt  Taking logarithm of both the sides,  log (ΔW) = (1/n).  log k + (1/n).  log t  (6)  (7)  The value of n can be obtained by the best-fit line of the slope for the log-log plots of (ΔW) against t. Further, the value of the general rate constant (k) can be determined by the intercept of the aforementioned best fit line to the y-axis. If the mass gain within a particular interval of exposure follows parabolic rate law, then the value of n in Eq. (6) would be 2, whereas a typical linear relationship would lead to a value of n1 [42,43]. The parabolic rate law implies that the oxidation reactions are controlled by diffusion of oxygen (atomic or anionic oxygen) from oxide-air interface to the oxide-composite interface through the protective oxide layer [44,45]. Assuming that the parabolic law has been followed in the current investigation, the value of parabolic rate constant (kp) can be determined by the following relation:  Table 1 Oxide scale thickness of composites after isothermal hold for 1 h, 8 h, and 24 h at 1300 °C.  Composition  Thickness of oxide scale formed after 1 h of exposure (μm) BSG + ZrO2  ZSBCL-7 ZSBCL-10 ZSBCL-14  57 ± 9 45 ± 2.3 86 ± 8  Thickness of oxide scale formed after 8 h of exposure (μm)  Thickness of oxide scale formed after 24 h of exposure (μm)  BSG  ZrO2  Porous ZrO2 + BSG  Total  BSG  ZrO2  Porous ZrO2 + BSG  Total  12 ± 2 17 ± 2.2 27 ± 4.3  11 ± 1 10 ± 1 12 ± 2.7  64 ± 3 55 ± 1.6 53 ± 3.2  87 ± 3.7 82 ± 2.9 92 ± 6  23 ± 4.3 26 ± 3 31 ± 1.7  13 ± 1.3 14 ± 1.5 13 ± 1.4  93 ± 3.7 73 ± 1.3 59 ± 6.5  129 ± 5.8 113 ± 3.6 103 ± 6.8  7  =()()()Kaaap**()feqLaOLaOBOLaBO,()2/214/74/2123232362=Kppap()*()()*()feqSiOSiOCOSiCO,()2apO2pO2pO2\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Table 2 The values of change in free energy per mole of oxygen (ΔGf), the equilibrium constant (Kf,eq), and the partial pressure of oxygen ( 1300 °C. Free energy of formation of oxidation products, ZrSiO4 and La2Si2O7 are also included.  ) for the oxidation reactions at  NO.  1 2 3 4  5 6  Reactions  (2/5)ZrB2 (s)+ O2 (g) = (2/5)ZrO2 (s) + (2/5)B2O3 (l) (2/3)SiC (s) + O2 (g) = (2/3)SiO2 (l) + (2/3)CO (g) (4/21)LaB6 (s) + O2 (g) = (2/21)La2O3 (s) + (4/7)B2O3 (l) SiC (s) + O2 (g) = SiO (g) + CO (g) Free energy of formation for ternary oxidation products. La2O3 + 2SiO2= La2Si2O7 ZrO2 + SiO2 = ZrSiO4  ΔGf (KJ/mol)  −566.9 −548.6 −132.9 −418.7  −2809.0 −1432.6  Kf,eq  6.7 × 1018 1.6 × 1018 2.6 × 104 8.0 × 1013  (atm)  1.5 × 10−19 6.0 × 10−19 3.8 × 10−5 1.2 × 10−14  Fig. 9. Typical plots of composites oxidized at 1300 °C showing the logarithmic plots of mass change (ΔW) against time for (a) 8 h and (b) 24 h; as well as plots of ΔW2 against time for (c) 8 h and (d) 24 h duration.  (ΔW)2 = kp.t + c  (8)  where kp is the parabolic rate constant, and c is a constant. Further,  d(ΔW)/dt = kp/2(ΔW)  (9)  The value of kp can be determined by the best-fit line of the slope for the plots of ΔW2 against t. The value of kp (= square of mass gain at unit time) is known to scale with the rate of mass gain by oxidation. Typical plots of log (ΔW) against log (t) for the isothermal duration of 8 h and 24 h are shown in Fig. 9(a) and (b), respectively, whereas those showing the variation of (ΔW)2 against t for the aforementioned duration are shown in Fig. 9(c) and (d). The value of oxidation exponent, general rate constant, and parabolic rate constant of the composites exposed at 1300 °C for 8 h and 24 h are presented in Tables 3 and 4, respectively. From the results in Table 3, the following inferences may be drawn: (i) the value of the general rate constant obtained for the ZSBCL-7 is an order of magnitude greater than that of either ZSBCL-10 or ZSBCL-14  Table 3 The values of oxidation exponent, the general rate constant, and the parabolic rate constant of composites oxidized at 1300 °C for 8 h.  Composite  ZSBCL-7 ZSBCL-10 ZSBCL-14  Time span (h)  1−8 1−8 1−5 5−8  Oxidation exponent (n)  General rate constant (k) (mgn cm−2n h-1)  2.8 2.1 1.9 3.4  60.56 8.1 8.2 118.7  R2  .98 .99 .99 .90  Parabolic rate constant (kp) (mg2 cm−4 h-1)  8.8 6.4 9.7 4.8  R2  .98 .99 .99 .84  composite; and (ii) the rate of oxidation of the ZSBCL-14 composite becomes slower (n = 3.4) after 5 h of isothermal exposure, which is consistent with the change in slope of the corresponding best-fit line in Fig. 9(a). Furthermore, considering that the values of n to be close to 2 for the initial stage of exposure (i.e. 8 h) for ZSBCL-7 or ZSBCL-10, and 5 h for the ZSBCL-14, the approximation involved in the use of  8  pO2pO2\\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Table 4 The values of oxidation exponent, the general rate constant, and the parabolic rate constant of composites oxidized at 1300 °C for 24 h.  Composite  Time span (h)  Oxidation exponent (n)  General rate constant (k) (mgn cm−2n h-1)  ZSBCL-7  ZSBCL-10  ZSBCL-14  1−24  1−24  1−8 8−24  3.9  3.0  2.9 9  486  45  45.6 9.9 × 106  R2  .99  .99  .99 .98  Time span (h)  Parabolic rate constant (kp) (mg2 cm−4 h-1)  1−8 8−24 1−8 8−24 1−8 8−24  7 3 5.4 2.8 5.6 0.9  R2  .99 .99 .99 .99 .97 .97  parabolic rate law as depicted in Fig. 9(c) and (d) appears quite appropriate. The results presented in Fig. 9(c) and listed in Table 3 show the value of kp as 9.7 mg2 cm−4 h-1 for the ZSBCL-14, which indeed is the highest among the investigated composites exposed at 1300 °C for 8 h, but is reduced to 5.5 mg2 cm−4 h-1 for the later stages of exposure, which is less than the values of kp for ZSBCL-7 (8.8 mg2 cm−4 h-1) and ZSBCL-10 (6.4 mg2 cm−4 h-1). Such a trend may also be considered as qualitatively consistent with the results of the mass gain against time (Fig. 2(c)). The values of n in the range of 3−4 for all three investigated composites exposed for 24 h at 1300 °C, as shown in Fig. 9(b) and Table 4, indicate that the average rate of oxidation during exposure for 24 h is significantly slower than that predicted by the parabolic rate law. Furthermore, the highest value of n (9) observed for the ZSBCL14 during the 8−24 h span of isothermal exposure indicates that its rate of mass change is significantly slower than that of ZSBCL-7 or ZSBCL-10, beyond the initial time-period of 8 h. It may be further noted that the plots in Fig. 9(d) represent the two different regimes of the rate of oxidation. In the first regime, which represents the early stage of oxidation (up to 8 h), the rate of oxidation of all the investigated composites is higher compared to that in the second regime (8-24 h). The higher rate of oxidation in the early stage of oxidation (≤8 h) may be attributed to the formation of an inadequately protective oxide scale during the initial stage of oxidation. Furthermore, in the second regime (8−24 h), the value of kp has been found to be similar for ZSBCL-7 and ZSBCL-10, whereas a significantly lower value is obtained for the ZSBCL-14 composite. Plots of the rate of mass change (dW/dt) calculated from Eq. (9) against time of isothermal exposure as shown in Fig. 10, indicate that (i) the rate of mass gain for the ZSBCL-14 composite is lower than that of ZSBCL-7 or ZSBCL-10 after the initial duration (9 h) of exposure, and the observed trend is similar to that observed for the kinetic parameters such as n and kp; (ii) for each of the investigated composites, a steady state characterized by a sharp decrease in the value of d(ΔW)/dt is reached beyond a limiting duration of 10−12 h; (iii) the transition in the slope of d(ΔW)/dt vs time plots is reached relatively earlier and the steady-state condition is most welldefined in the ZSBCL-14 composite; and (iv) at the end of 24 h, the values of d(ΔW)/dt decrease in the following order, ZSBCL-7 > ZSBCL10 > ZSBCL-14. The commonly observed decrease in the rate of oxidation with the increase in the duration of exposure from 8 h to 24 h at 1300 °C for each of the investigated composites may be attributed to the progressive evolution of a compact and more protective oxide scale, which acts as an effective diffusion barrier for oxygen. According to the available literature, the value of n is reported to be 2 for SiC containing UHTCs at an intermediate temperature range (1300−1600 °C), and this has been rationalized by considering the diffusion of oxygen through the viscous BSG layer as the rate-limiting step [16,46]. However, it deviates from the parabolic rate law to a higher-order, if the rate of diffusion of oxygen through the protective oxide scale is further reduced [47]. Guo et al. [23] have reported that increasing the amount of SiC from 10 to 30 vol% in the ZrB2-SiC composite, changes the oxidation kinetics from parabolic to cubic due to the formation of a higher amount of SiO2 in the oxide scale. In the current study, the growth of BSG layer also plays a similar important role in oxidation kinetics. After forming an appreciable amount of BSG  layer, the oxidation exponent has been found to be 9 for the ZSBCL14 composite at 1300 °C for 24 h exposure. Furthermore, the value of kp (5−10 mg2 cm−4 h-1) obtained after 8 h of exposure in the present study is in good agreement with the results reported by Kakroudi et al. [48] on the ZrB2-SiC-TaC composites, which has shown the value of kp to vary from 2.2-30.6 mg2 cm−4 h-1 with increase in the temperature of isothermal exposure from 1000 °C to 1400 °C for 10 h duration. ZapataSolvas et al. [25] have reported that the addition of La2O3 drastically enhances the oxidation resistance with n 8 at 1600 °C by forming a protective MeOxCy (Me = Zr, Hf, or Si) phase. These results are comparable with the current study showing the value of n  9 for the ZSBCL-14 composite after forming a sufficiently thick protective oxide scale at 1300 °C. Obviously, the kinetics of oxidation at 1300 °C in the present study is expected to be slower than that observed in the earlier investigation at 1600 °C [25].  4.3. Evolution of oxide scale and its relationship with kinetic parameters  evolving  4.3.1. Effect of oxidation products on structure and integrity of BSG layer The examination of the oxide scales formed on the investigated composites after isothermal exposure for various durations at 1300 °C through XRD and SEM-EDS analyses has revealed that La2Si2O7, SiO2, ZrSiO4, and ZrO2 are the main products of oxidation. The positive component of mass change observed during oxidation is ascribed to the formation of non-volatile products including ZrO2, SiO2, La2Si2O7, and ZrSiO4, whereas the negative component can be explained on the basis of the formation of gaseous products like CO2 (g), CO (g), and B2O3 (g) followed by their escape, as well as occasional spallation of a part of the oxide scale triggered by growth-related internal stresses. It is necessary to understand the effect of various constituent phases on the formation of a protective glassy scale, which may vary with the net composition. It is well-known that increase in the amount of B2O3 in the borosilicate glass increases its fluidity and plasticity with lowering of viscosity [49]. On the other hand, its vaporization causing increase in the proportion of SiO2 enhances the oxide scale viscosity. Further, earlier studies have shown that the presence of Zr2+ ions in borosilicate glass acts as a glass network modifier and increase the glass-forming ability. The SiO4 − tetrahedral network is strengthened by lowering the fraction of non-bridging oxygen atoms created by vaporization of B2O3 due to the connections established between the ions of (SiO4)and (ZrO6)2[50,51]. It has been proposed in an earlier study that as a result of increase in viscosity due to the dissolution of a significant amount of ZrO2, the molten BSG is prevented from ascending to the top surface from the ZrO2+BSG mixed layer [27]. Further, mixing of a significant amount of undissolved ZrO2 inclusions (beyond percolation threshold) with BSG may be detrimental, as the former oxide provides easier diffusion path for oxygen anions. It has been shown that addition of a small amount of La2O3 (1 mol %) to B2O3-SiO2 (binary) system causes strong phase separation due to miscibility gaps in binary SiO2-La2O3 and B2O3-La2O3 systems [52]. In a similar manner, mixtures having higher B2O3 content (40−60 mol%) with moderate SiO2 content (20−40 mol%) have exhibited separation of the La-rich phase. On the contrary, addition of a large amount of La2O3 (40−60 mol%) to B2O3-SiO2 glass leads to crystallization with  9  \\x0c', 'F  i  g  .  0 1  .  P  l  o  t  s  d  e  p  i  c  t  i  n  g  t  h  e  v  a  r  i  a  t  i  o  n  i  n  r  a  t  e  o  f  m  a  s s  c  h  a  n  g  e  a  g  a  i  n  s  t  t  i  m  e  f  o  r  c  o  m  p  o  s  i  t  e  s  e  x  p  o  s  e  d  i  s  o  t  h  e  r  m  a  l l  y  a  t  0 0 3 1  °  C  f  o  r  4 2  h  .  S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  10  \\x0c', 'S.K. Kashyap, et al.  de-polymerization along with an increase in the fraction of non-bridging oxygen anions, which in turn lowers the glass-transition temperature. In such cases, the La3+ ions in La2O3 play the role of chargecompensation by occupying positions close to borosilicate network and transferring non-bridging oxygen anions to SiO4 − and BO3 units [53]. On the other hand, mixtures with almost equal amounts of SiO2 (40−60 mol%) and B2O3 (30−60 mol%) with 20−30 mol% La2O3 have formed homogeneous glass, where the La3+ ions have bonded strongly to non-bridging oxygen anions and thereby acted as glassnetwork modifier, leading to increase in the glass transition temperature. It has been earlier reported that the SiO2 tetrahedron with La atoms in interstitial positions is more stable than that having those at substitutional sites, because the interstitial La breaks the nearest Si-O bond to form La-O and La-Si bonds, which is beneficial for improving the high‐temperature stability of SiO2 [54]. In another study, it has been demonstrated that the formation of La2O3 near SiC promotes its inclusion into the BSG phase, which in turn increases both its viscosity and thermal stability, and thereby lowers the oxygen diffusivity [25]. Further, phase separation with La2O3 addition to the BSG due to miscibility gap is expected to increase the viscosity and lower oxygen diffusivity [27]. Based on the calculated isothermal section of the ZrO2-SiO2-B2O3 phase diagram at 1500 °C, increase in the B2O3 content of the BSG liquid enhances its solubility for ZrO2 resulting in the formation of ZrO2-SiO2-B2O3 (BSZ) liquid. In a similar manner, La2O3 is also expected to dissolve in the BSG liquid. However, vaporization of B2O3 is expected to cause enrichment of ZrO2 and La2O3 in the oxide scale, leading to formation of ZrSiO4 and La2Si2O7, respectively [23,27], whose presence has been confirmed in the present study by both XRD analysis and EDS examination of the phases present in the oxide scales, as shown in Figs. 3-6.  4.3.2. Effect of LaB6 content on protective character of oxide scale In the present study, the microstructural study of top surface and cross-section of the oxide scale has revealed the absence of a continuous layer of oxide scale after 1 h of isothermal exposure [Figs. 4,7 and 8(a)], which is in tune with the random fluctuations in mass change from positive to negative during the initial stage of oxidation. The oxide scale formed after exposure at 1300 °C for 1 h, is expected to contain relatively less amount of SiO2 compared to that of B2O3 due to faster oxidation of both LaB6 and ZrB2 phases. Partial vaporization of B2O3 may along with spallation have contributed to randomly observed decrease in mass, as shown in Fig. 2(a). Interestingly, in the course of isothermal exposure for 1 h, the oxide scale developed on the ZSBCL-7 composite shows the presence of open pores. On the other hand, the ZSBCL-10 and ZSBCL-14 composites show relatively insignificant cracking on the surface of the oxide scale. This observation could be attributed to the formation of a greater amount of B2O3 due to the presence of much higher amount of LaB6 in both ZSBCL-10 and ZSBCL-14 composites. It may be noted that the amount of B2O3 formed by oxidation of a mole of LaB6 is expected to be 3 times more than that by a similar amount of ZrB2. The formation of a larger amount of B2O3 is expected to increase the fluidity of the BSG scale, and facilitate self-healing of cracks and porosities. However, due to the formation of a relatively smaller amount of La2O3 compared to that of B2O3 (being contributed by oxidation of both LaB6 and ZrB2), phase separation within the BSG may be responsible for identification of the former phase in the XRD patterns from the oxide scales formed through exposure for 1 h at 1300 °C (Fig. 3(a)). At the locations having a larger amount of La2O3 formed along with BSG containing nearly equivalent amounts of B2O3 and SiO2, crystallization with formation of La2Si2O7 may have happened, and therefore this phase appears to be quiet scattered in the resultant oxide scale (Fig. 4). As mentioned above, a limited dissolution of ZrO2 in the BSG during the first hour of exposure is expected to have strengthened the glass with lowering the possibility of crystallization. A few earlier reports have shown that oxidation kinetics depends to  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  a large extent on the presence of open pores on the oxide scale surfaces, formation of cracks and spallation of oxide scale, porous structure beneath the BSG layer, and SiC-depleted region [55,56]. Therefore, it may be speculated that due to the presence of the highest amount of LaB6 in the ZSBCL-14 composite, creation of pores in the oxide scale by vaporization of B2O3 has probably enhanced the oxidation rate during the initial stages (0−5 h) of exposure at 1300 °C, compared to that observed in other two composites. On the other hand, during the time span of 8−24 h at 1300 °C, the amount of SiO2 formation appears to have increased proportionately to the other constituents of the oxide scale, and this may have accelerated the formation of a homogenous BSG layer within the oxide scale, which is considered as primarily responsible for protection against further oxidation. As the amount of La2O3 (having more mass than an equivalent amount of ZrO2 formed by oxidation of ZrB2) and B2O3 formed in the oxide scales of the investigated composites with less LaB6 content is expected to be less too, the ZSBCL-7 appears to have exhibited the highest oxidation resistance during the span of 0−8 h. Although the relative density of the ZSBCL-7 (98 %) is slightly less than that of ZSBCL-10 (99 %) and ZSBCL-14 (99.5 %), this factor too is not expected to have a significant effect on kinetics of oxide scale formation. The average sizes of ZrB2 grains in the microstructures of the investigated composites have been found to have only minor variations (2.6−3.7 μm), depending on the amount of LaB6 content, which suggests that this may not have affected the kinetics of oxide scale growth significantly. Therefore, the role of ZrB2 grain boundaries and ZrB2-SiC interfaces appears to be similar for the investigated composites, and may not have caused much difference in the oxidation behavior.  4.3.3. Effect of LaB6 content on BSG layer growth Investigation of the oxide scales developed after 8 h or 24 h of isothermal exposure shows the well-developed and distinct layers of BSG, and La2Si2O7 on the outermost surface of typically multi-stratified oxide scales [Figs. 5, 6 and 8(b-c)]. The thickness of the BSG layer formed after 8 h of isothermal hold at 1300 °C as shown in Table 1, is found to be increased from 12 μm (in ZSBCL-7) to 32 μm (in ZSBCL-14) with increasing volume fraction of LaB6, while the ZrO2 layer thickness for all three composites is found to be almost similar (10−13 μm), irrespective of the LaB6 content. Such an increase observed in the BSG content alone with the amount of ZrO2 remaining unchanged with increasing amount of LaB6 in the composites suggests that the SiC particles at sub-surface locations are preferentially oxidized. Such preferential oxidation of sub-surface located SiC particles may be ascribed to lower partial pressure of oxygen required for oxidation of SiC compared to that for ZrB2, as shown in Table 2. Also, there is a strong possibility of active oxidation of a part of SiC through Reaction (4) in Table 2, and the escaping SiO (g) is probably converted to SiO2 (s) at the oxide-air interface. The escape of SiO(g) and CO(g) from the reaction front is expected to further enhance the oxidation rate of SiC, as is normally expected according to the law of mass action. Interestingly, the formation of porous ZrO2 with traces of BSG due to the ingress of oxygen from the outer scale appears to be partially restricted with the increasing volume fraction of LaB6. Measurements of thickness of various layers within the oxide scales after exposure at 1300 °C for 24 h have shown that the porous ZrO2 + BSG is 64 ± 3 μm thick in case of the ZSBCL-7 composite, whereas it is only 53 ± 2.2 μm thick in case of the ZSBCL-14, indicating that the oxide scale is more protective in case of the latter composite. Based on the measurements carried out using the SEM images and the data presented in Table 1, plots depicting the variation of the intermediate BSG layer thickness with LaB6 content during the time-spans of 1−8 h and 8−24 h are presented in Fig. 11(a) and (b), respectively. According to the results shown in Fig. 11(a), the growth rate of the BSG layer within the oxide scale at the end of 8 h isothermal exposure is found to scale with increasing LaB6 content of the investigated composites following a linear  11  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 11. Plots depicting the variation of (a) BSG layer growth up to 8 h with LaB6 content; and (b) BSG layer growth during 8-24 h span with LaB6 content. Best-fit lines are drawn in both the plots.  relationship, t(BSG) = 7.5x + 3.66 with R2 = 0.96. This observation may be ascribed to the increased formation of B2O3 by oxidation of LaB6, which probably has lowered the viscosity of the BSG layer during isothermal exposure at 1300 °C, and this in turn may have enhanced the diffusivity of oxygen, leading to increase in the growth rate. However, the growth rate of the BSG layer appears to decrease with increasing LaB6 content during the span of 8−24 h following a linear relationship, as evidenced by the linear-fit of data, t(BSG) = -3.5x + 15 with R2 = 0.94 in the plot shown in Fig. 11(b). For example, the thickness of BSG layer has increased by 91 % or 54 %, respectively for ZSBCL-7 or ZSBCL-10 composites on increase of the isothermal exposure duration from 8 h to 24 h, but such increment (15 %) is significantly less in ZSBCL-14 composite. Further, by comparing the results in Fig. 11(a) and (b), it is obvious that for each of the investigated composites with given LaB6 content, the observed growth rate during the span of 8−24 h is significantly less than that recorded for the duration up to 8 h. This observation also suggests that there is a reversal in the growth rate of the BSG layer, after a critical thickness is reached for a given composite. In order to explain this observation, it is necessary to consider that with increasing duration of exposure, the amount of B2O3 dissolved in the BSG layer may have decreased by vaporization, and this resulted in increased fractions of dissolved ZrO2, La2O3, and SiO2 as a natural consequence, causing a simultaneous increase in the viscosity. It is well-expected that an increase in the viscosity of the BSG layer would significantly lower the diffusivity of oxygen. Such an increase in the volume fractions of the aforementioned non-volatile oxides may therefore be considered as the possible explanation for the observation of lower rate of growth in the BSG layer during the 8−24 h span compared to that happening on exposure within 1−8 h time-period for each of the investigated composites, irrespective of the LaB6 content. A similar reason may also be proposed to explain the observed decrease in the growth of BSG layer thickness during the 8−24 h exposure span with increase in the LaB6 content of the investigated composites.  4.3.4. Effect of BSG layer growth on oxidation kinetics After a continuous BSG layer is formed, it is appropriate to infer that diffusion of oxygen through the protective BSG layer is the rate limiting factor in case of the investigated composites, because the diffusivity of oxygen through SiO2 (10−21 m2/s at 1550 °C) is slower by several orders of magnitude than that in ZrO2 (10−10 m2/s at 1500 °C) [57,58]. In order to correlate the oxidation kinetics with growth of the oxide scales of the investigated composites, the values of parabolic rate constant, kp are plotted against the BSG layer thickness recorded after exposure for 8 h and 24 h, as shown in Fig. 12(a) and (b), respectively. As shown in both Fig. 12(a) and (b), the values of parabolic rate constant, kp decreases with increasing thickness of the BSG layer, whereas a linear relationship appears to be followed during exposure up to 8 h, such that the data fall in a single best-fit line (kp = -0.25*t(BSG) + 11.36) with R2 value of 0.91. As the oxidation rate depends on the diffusion of oxygen through the BSG layer, it is intuitive that a thicker layer of BSG would lead to increase in the diffusion distance, and therefore a smaller value of kp, as shown in Fig. 12(a) and (b). The  absence of a single best-fit line for the data in Fig. 12(b) may be ascribed to the difference between the compositions of the BSG layer formed in the oxide scales of the investigated composites, which in turn is expected to have altered the oxygen diffusivities through these layers. It may be noted that in Fig. 12(a), the selected kp values represent a span of 0−8 h for ZSBCL-7 and ZSBCL-10, whereas it is obtained for the span of 5−8 h in case of the ZSBCL-14 composite. The choice of kp representing the span of 5−8 h in case of the ZSBCL-14 composite has been made on the basis of the following considerations: (i) the oxidation resistance of the investigated composites is primarily facilitated by the formation of BSG layer within the oxide scale, and (ii) the kp value calculated in case of the ZSBCL-14 composite has decreased sharply on moving from the span of 0−5 h to 5−8 h as shown in Fig. 9(c), which indicates that the BSG has grown in thickness, or its composition has changed more drastically during the latter time-span, such that the ingress of oxygen is significantly retarded. On the other hand, the kp values observed in case of the ZSBCL-7 and ZSBCL-10 composites have remained more or less constant during the span of 0−8 h, suggesting that the BSG layer growth rate along with change in viscosity through compositional changes may have taken place somewhat uniformly throughout this time-period, except for the first hour, after which the presence of this layer could not be detected through microstructural examinations (Table 1). Further, the observed decrease in the values of the rate of mass change, d(ΔW/dt) with time (as shown in Fig. 10) after the first hour of exposure may be considered as an indirect confirmation of a continuous evolution of the protective BSG layer. As inferred from the trend in Fig. 10 and the kinetic data in Tables 3 and 4, the ZSBCL-7 has exhibited the lowest rate of mass gain during the initial 8 h of exposure, which is followed by ZSBCL-10 and ZSBCL14 composites, whereas the trend is completely reversed after an exposure duration of 8−9 h. This observation may be ascribed to faster rate of oxidation of LaB6 compared to that of ZrB2 and SiC, and therefore the rate of mass gain by oxidation in the initial stages appears to scale with the LaB6 content of the investigated composites. Once the LaB6 present near surface is almost fully converted to La2O3, which probably takes 8−9 h, the oxidation process is controlled primarily by oxidation of SiC with formation of BSG and La2Si2O7 along with increase in viscosity (as discussed in Section 4.1 and earlier part of this section), thereby enhancing the protective character of the oxide scales. Considering that the mass gain by the ZSBCL-14 is more than that of the ZSBCL-10 up to the exposure duration for 8 h, it is inferred that a critical minimum duration is required for a desirable ratio of oxide scale constituents to form, so that a protective glassy layer with desirable amount of viscosity is formed within the oxide scale. This hypothesis is in line with the observations recorded from Fig. 12(a), indicating that the faster the growth rate of the BSG layer in the oxide scales, the lower is the observed value of kp from the isothermal oxidation tests on the investigated composites. This observation is justified, considering the increase in diffusion distance of oxygen with increasing growth of the BSG layer within the oxide scale. If the time-span of isothermal exposure up to 8 h at 1300 °C is considered for the ZSBCL-14 composite, a transition in the mass gain kinetics is observed after the initial period of  12  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  Fig. 12. Plots showing the change in the values of kp with thickness of BSG layer after isothermal exposure of (a) 8 h, and (c) 24 h duration.  5 h, as is obvious from the results shown in Table 3. Furthermore, on crossing over from the time span of 0−8 h to 8-24 h, the value of oxidation exponent (n) calculated for the ZSBCL-14 composite is increased very sharply from 2.9 to 9, whereas the parabolic rate constant (kp) is decreased from 5.6 to 0.9 mg2 cm−4 h-1. Such an observation regarding sharp changes in oxidation rate (kinetics-related) parameters is consistent with the trend showing lower rate of mass gain in the ZSBCL-14 compared to that in ZSBCL-7 or ZSBC-10 (Fig. 10), which in turn appears to be related to significantly reduced rate of the BSG layer growth in the oxide scale of the ZSBCL-14 during the span of 8−24 h. Such an observation may be explained by the requirement for the formation of a critical thickness of dense and continuous BSG layer for inhibiting the oxygen penetration and the consequent enhancement of the oxidation resistance of the composites. Obviously, the present observations suggest that such a critical BSG layer thickness is reached earliest in the oxide scale of the ZSBCL-14 composite, which in turn leads to a sharp decrease in kp after exposure for 5 h at 1300 °C.  and B2O3 formed by its oxidation, along with the continuous change in the amount of the latter constituent due to vaporization during exposure at 1300 °C. The results of the present study have significantly improved the scientific understanding of the stages of oxide scale evolution with formation of the protective BSG layer, and its strong dependence on the LaB6 content during exposure at 1300 °C. One may prefer to add LaB6 rather than La2O3 to the investigated ZrB2-SiC composites, considering that additional amount of B2O3 formed may improve the self-healing ability of the BSG layer developing within the oxide scale by increasing its plasticity and fluidity by decreasing the viscosity. At the same time, it is necessary to limit the decrease in viscosity by proper control of B2O3 content, so that inward diffusivity of oxygen is not increased significantly. An understanding of the effect of LaB6 addition on oxidation resistance of the investigated ZrB2-SiC composites may be considered as a stepping stone for further development of this material in potential high temperature applications involving components in aero-engines and space vehicles.  4.4.  Impact of  the present study  The present study has led to a number of interesting inferences regarding oxidation kinetics and evolution of oxide scale in the investigated ZrB2-SiC based UHTCs having varying LaB6 content, besides generating useful data for isothermal exposure at 1300 °C. Analysis of the results has shown the existence of a direct correlation between the oxidation kinetics and the growth rate of the BSG layer, which primarily comprises the protective part of the oxide scale. As the diffusion of oxygen through the BSG layer is the slowest process in the sequence of events involved in oxidation of the investigated composites, changes in its thickness and viscosity have directly influenced the kinetics of their mass gain. Interestingly, linear relations with slopes having opposite signs have been found for suitably expressing the dependence of the BSG growth rate with LaB6 content for durations of 1−8 h and 8−24 h, Fig. 11. Moreover, for a given composition of the LaB6 containing composite, it takes some time of exposure at high temperature before the protective character is fully established, which coincides with a critical thickness, as well an optimum BSG layer composition and viscosity being reached. As the protective BSG layer appears to be developing continuously in the oxide scale after about an hour of exposure, the oxidation kinetics is slowed until a steady state is reached after duration of 10−12 h. The present study has shown that on exposure at 1300 °C, the least mass gain is observed in the ZSBCL-10 composite after exposure for 8 h, and in the ZSBCL-14 after 24 h, which suggests that the LaB6 content has a significant effect on the kinetics of oxidation and protective scale formation. It has been further inferred that the major cause for change in oxidation behavior with the LaB6 is of course the variation in oxide scale chemistry and resultant viscosity due to varying amounts of La2O3  13  5. Conclusions  The isothermal oxidation behavior of ZrB2-SiC-LaB6 composites prepared by SPS at 1600 °C has been examined at 1300 °C for 1 h, 8 h or 24 h. Based on the results obtained in the present investigation, the following conclusions can be drawn:  1) The oxidation tests carried out at 1300 °C for 1 h duration have shown random fluctuations in mass change, because a stable and protective BSG layer is not able to develop during the early stage of oxidation. 2) The formation of an appreciable amount of BSG along with a continuous layer of La2Si2O7 on the outermost surface hinders the diffusion of oxygen through the oxide scale; and therefore the mass gain follows approximately parabolic rate law during exposure up to 8 h; whereas it deviates from parabolic to higher-order kinetics for a longer duration of exposure at 1300 °C. 3) The parabolic rate constant is found to scale with the BSG layer thickness, as the diffusion distance for oxygen to reach the oxide scale-composite interface is increased. Decrease in the rate of mass gain with time right after the first hour is suggestive of the growth of the protective BSG layer during this time. 4) The oxidation of ZSBCL-14 has exhibited faster kinetics than the other two investigated composites up to 8 h of exposure, whereas after this period, a significant drop in the rate of oxidation is observed due to the formation of a compact BSG scale with optimum thickness and viscosity, as well as a continuous La2Si2O7 layer on the outermost surface, both of which may have jointly hindered further ingress of oxygen. 5) Microstructural analysis  cross-sections  scale  has  of  the  oxide  \\x0c', 'S.K. Kashyap, et al.  Journal of the European Ceramic Society xxx (xxxx) xxx-xxx  revealed that with an increase in the amount of LaB6, the total thickness of BSG layer reached at the end of 8 h or 24 h is increased, whereas that of the porous ZrO2+BSG layer is reduced. However, the growth rate of the BSG layer during the 8−24 h span of exposure is reduced with respect to that during the 1−8 h time-period, which is suggestive of its protective nature evolving in this period. Both during the initial 1−8 h and later 8−24 h exposures, the growth rate of BSG has shown linear correlation with the amount of LaB6 in the composite, with the slopes of the best-fit lines in these two regimes exhibiting opposite signs.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgements  Technical assistance received from Mr. Mithun Das and Mr. B. Santu Mudliyar, Staff members of Central Research Facility, IIT Kharagpur for characterization of specimens, is gratefully acknowledged.  References  [3]  [6]  [1]  S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. 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},{
  "_id": 118,
  "PDF": "Long time ablation behaviors of designed ZrC-SiC-TiC ternary coatings for environments above 2000 °C.pdf",
  "Text": "['Corrosion Science 170 (2020) 108645  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Long time ablation behaviors of designed ZrC-SiC-TiC ternary coatings for environments above 2000 °C  T  Xiaohui Pana,b, Yaran Niua,*, Xueting Xua, Xin Zhonga, Minhao Shia, Xuebin Zhenga,*, Chuanxian Dinga  a Key Laboratory of Inorganic Coating Materials CAS, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, 100049, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: ZrC-SiC-TiC ternary coatings Ablation resistance Oxide products Vacuum plasma spray  Thermal protection systems with long time anti-ablation property for applications in ultrahigh temperature environments have been a choke point for the development of hypersonic vehicles. In the present work, new ZrCSiC-TiC ternary composite coatings were designed and fabricated by vacuum plasma spray. The ablation behaviors of the ternary coatings were evaluated by a plasma flame with temperatures about 2200−2500 °C and compared with the ZrC-SiC and ZrC-TiC composite coatings. The characteristics of phases, element distributions and microstructure of the oxidized layers were observed in detail. The results showed that the anti-ablation resistant property of the ternary coatings was much better than those of the ZrC-SiC and ZrC-TiC composite coatings. The oxidation products, like high viscous SiO2 and high thermo stable TiO2, played a mutual promoted role in modifying the oxide layers, which contributing a lot to the improved ablation resistance. The ablation mechanisms of the ZrC-SiC-TiC ternary coatings were analyzed based on microstructure observations and thermodynamic analysis.  1.  Introduction  The severe operating conditions of hypersonic aircrafts such as ultra-high temperature, strong airflow scour and high thermal stresses have provoked urgent requirements for novel thermal protection systems (TPS) that have good oxidation resistance, thermal shock resistance and ablation resistance [1-4]. Ultra-high temperature ceramics (UHTCs) exhibit striking multi-characters, including extremely high melting points, high strength and modulus, relatively low density and so on, which are potential candidate materials for TPS [5-10]. ZrC has relatively low density (6.73 g/cm3), achigh melting point (3540 °C), ceptable resistance in ablation environment and low cost, which makes it attractive among all UHTCs. However, the application of mono ZrC is limited by its poor oxidation resistance. ZrC is susceptible to oxidation when temperature rises above 600 °C, forming porous ZrO2 scale that provide channels for oxygen diffusion [11-15]. Over the past years, SiC has been a major choice to improve the anti-ablation property of ZrC [16-21]. Jia et al. [16] fabricated ZrC20 vol.%SiC coating on C/C composites and investigated its ablation behavior usiing oxyacetylene flame. They found that after 120 s ablation, the linear ablation rate reduced by 96.4 % and the mass gain rate  increased by 383.3 % compared with those of the pure ZrC coating. The reports of ablation behaviors of ZrC-SiC systems were summarized in Table 1. It can be concluded that ZrC-SiC exhibited good ablation resistance in environments below 2200 °C for short time. The incorporation of SiC into ZrC could form a protective ZrO2-SiO2 oxide scale. SiO2 could serve as an ideal oxygen diffusion barrier because of its low for example, 3.2 × 10−15 gcm−1s−1 at 1000 °C oxygen permeability, [22]. However, the effect of SiC is limited by the environmental temperature for the following reasons: (i) the high evaporation and decomposition trend; (ii) the active oxidation of SiC inside. The evaporation or decomposition of SiO2 would leave a porous oxide scale, providing channels for oxygen diffusion. Besides, the active oxidation of SiC would generate large amount of gaseous products and cause damages. In the previous work in our laboratory, we designed and fabricated ZrC-TiC composite coatings with different TiC contents (20, 30, 40 vol. %) using vacuum plasma spray [23]. The ablation behaviors of this kind of coating were investigated under a plasma flame. The results showed that the ZrC-30 vol.%TiC coating exhibited better ablation resistance than that of the ZrC-30 vol.%SiC coating after 150 s ablation. The ZrC30 vol.%TiC coating kept integrated while the ablation center of ZrC ⁎ Corresponding authors. E-mail addresses: yrniu@mail.sic.ac.cn (Y. Niu), xbzheng@mail.sic.ac.cn (X. Zheng).  https://doi.org/10.1016/j.corsci.2020.108645 Received 6 September 2019; Received in revised form 23 March 2020; Accepted 25 March 2020  Available online 04 April 2020 0010-938X/ © 2020 Elsevier Ltd. All rights reserved.  \\x0c', 'X. Pan, et al.  Table 1  Summary of ablation behaviors of ZrC-SiC systems.  System  ZrC-20 vol.%SiC ZrC-20 vol.%SiC ZrC-SiC ZrC-SiC ZrC-30 vol.%SiC ZrC-SiC ZrC-SiC  System type  Ablation time  Ablation temperature  Coating Coating Coating Coating Ceramic Coating Coating  120 s 120 s 5−45 s 60 s 600 s 60 s 20 s  1738 °C 2105 °C 2173 °C -  2100 °C 2100 °C -  Corrosion Science 170 (2020) 108645  Results  Integrated Damaged Integrated Integrated Damaged Integrated Integrated  Refs  [16] [16] [17] [18] [19] [20] [21]  30 vol.%SiC peeled off. Based on the microstructure and thermal dynamic analysis, it was found that (Zr, Ti)O2 and TiO2 are more thermodynamically stable than SiO2, which contributed to the improved ablation resistance. However, we also found that the ZrC-30 vol.%TiC coating was almost totally oxidized after 150 s ablation. This phenomenon was attributed to the relatively high oxygen diffusion coefficients (1.12 × 10−13 m2s−1, of both TiO2 at 1800 °C) and ZrO2 (1.16 × 10−12 m2s−1, at 1800 °C) [24]. Things could become interesting if the advantages of SiO2 (low oxygen permeability) and TiO2 (high thermo stability) could be combined. However, the related works have not been reported yet. In the present work, ternary ZrC-SiC-TiC coatings were designed and prepared by vacuum plasma spray. The ablation behavior of two kinds of ternary composite coatings (ZrC-20 vol.%SiC-10 vol.%TiC, ZrC10 vol.%SiC-20 vol.%TiC) were evaluated by a plasma flame at atmosphere. ZrC-SiC (30 vol.%) and ZrC-TiC (30 vol.%) composite coatings were also prepared for comparison. It is expected that the SiC and TiC additions could play a positive joint function on the anti-ablation property of the ZrC coating.  2. Experimental  2.1. Coating preparation and characterization  Powder mixtures with different compositions, including ZrC-30 vol. %SiC (denoted ZS3), ZrC-20 vol.%SiC-10 vol.%TiC (denoted ZS2T1), ZrC-10 vol.%SiC-20 vol.%TiC (denoted ZS1T2) and ZrC-30 vol.%TiC (denoted ZT3), were prepared by mixing ZrC (Changsha Langfeng Metal Materials Co., Ltd., China), SiC (Qinhuangdao Yinuo Advanced Material Co., Ltd., China) and TiC (Qinhuangdao Yinuo Advanced Material Co., Ltd., China) by ball mixing in alcohol. Then, the powder mixtures were agglomerated by spray drying and sintered to obtain powders that were (Φ 40 mm × 5 mm) suitable for vacuum plasma spray. Graphite (Beijing Jinglong Carbon Materials Co., Ltd., China) with a SiC bonding layer was used as substrates. The preparation details of the SiC bonding layer could be seen elsewhere [25]. Vacuum plasma spray (A-2000, Sulzer Metco AG, Switzerland) was used to fabricate the coatings.  Hydrogen and argon were used to generate the plasma flame. The details of coating preparation could be seen in our previous work [26].  2.2. Ablation test  Argon and hydrogen were used to generate the plasma flame with a heat flux of 3.01 MW/m2 for evaluating the ablation resistance of the composite coatings. The heat flux was measured by a Gardon heat flux meter (Shanghai Tuxin electronic technology co., Ltd., China). All the ablation resistance tests were conducted at atmosphere for 300 s and 600 s, respectively. It should be stated that there was a certain time (about 30 s) for the plasma flame generator to cool down after a duration of 300 s during the 600 s ablation test. The coatings after 300 s and 600 s ablation were labeled as, taking the ZS3 coating for example, ZS3-300, ZS3-600, respectively. The temperatures of the ablation centers of the coating samples were recorded by a two-colour infrared pyrometer (Marathon MR1SC, Raytek, USA).  2.3. Composition and microstructure characterization  The phase compositions of the as-sprayed and ablated coatings were characterized by X-ray diffraction with Cu Ka (λ → 1.54056 Å) radiation (XRD, RAX-10, Rigaku, Japan). The step size and scanning rate are 0.005°and 10°/min, respectively. The microstructures and chemical element distributions of the as-sprayed and ablated coatings were characterized using a field emission scanning electron microscope (FEISEM, Magellen 400, USA) equipped with an energy dispersive spectrometer (EDS, PN-5502, INCA ENGERY, UK). The ablation centers of the coatings were chosen to evaluate the ablation resistance and take detailed microstructrue characterizations. The particle size of powders was evaluated by a laser particle analyzer (BT-9300 s, China). The porosity of the as-sprayed coatings was evaluated by three cross-sectional images with a magnification of 1000× using an image analysis software (Leica Qwin, Germany).  Fig. 1. Morphologies and particle size distribution of ZS3 starting powders.  2  \\x0c', 'X. Pan, et al.  Fig. 2. XRD patterns of as-sprayed ZS3, ZS2T1, ZS1T2 and ZT3 coatings.  3. Results  3.1. Phase compositions and microstructures of as-sprayed coatings  Fig. 1 shows the typical morphologies and particle size distributions of the ZS3 feedstock powders. It can be seen that the powder was (D50) was about 35−40 μm spherical and the measured mean size (Fig. 1b). Fig. 2 displays the XRD patterns of the as-sprayed coatings. The coatings were mainly composed of cubic ZrC. SiC could be observed in the ZS3 coating, but not in the ZS2T1 and ZS1T2 coatings, which could be ascribed to the relatively low content of SiC for those coatings. The positions of ZrC peaks were shifted about 0.5° higher, for the solid solution phenomenon of ZrC and TiC [27]. Fig. 3 shows the surface and cross-sectional morphologies of the as-sprayed coatings. Fully molten and partially molten particles could been seen in Fig. 3(a-d). No obviou cracks or voids could be observed between the SiC bonding layer and the top coating, which indicated good bonding for the coatings (Fig. 3(e-h)). The thickness of the ZS3, ZS2T1, ZS1T2 and ZT3 coatings were about 310, 360, 390 and 420 μm, respectively. The porosity of the composite coating was calculated using the crosssection SEM images, which were about 5-8 %.  3.2. Ablation behavior  Fig. 4 shows the macrographs of the coatings after 300 s and 600 s ablation tests. It can be seen that the morhologies of all coatings changed after the ablation tests and the ablation time had a significant influence. Only part proportion (about 40 %) of ZS oxide scale was retained after 300 s ablation (Fig. 4a), while the ZS2T1, ZS1T2 and ZT3  Corrosion Science 170 (2020) 108645  (Fig. 4b-d) coatings kept integrated (100 %) and seemed still bonded well to the substrate. When the ablation time was extended to 600 s, the unpeeling-off part of ZS coating is further decreased to 10 %, as shown in Fig. 4e (Yellow circle area). With the increasing of the TiC content, the ablated coating become more integrated. However, the ZT3 coating peeled off after 600 s ablation (Fig. 4h). The temperature/time curves of the coating samples during the ablation tests are illustrated in Fig. 5. The final temperature was decreased with the increasing of the TiC content, which were 2390 °C, 2363 °C, 2300 °C and 2266 °C, for the ZS, ZS2T1, ZS1T2 and ZT3 coating, respectively. The XRD patterns of the ablated coatings are shown in Fig. 6. For the coatings after 300 s ablation, it can been seen that dominant phase of all the coatings was monoclinic zirconia (m-ZrO2) (Fig. 6a). A trace of (Zr, Ti)O2 could be found in the ablated ZT3 coatings. (Zr, Ti)O2 was generated through reaction 1.  ZrO2(s) + TiO2(l) → (Zr, Ti)O2(l)  (1)  After 600 s ablation, phase transformation could be figured in Fig. 6b. That is, the intensity of m-ZrO2 dramatically declined and the peaks of (Zr, Ti)O2 became more distinct in all the TiC contained coatings compared with the 300 s ablated coatings. SiC could be found for the ZT3-600 coating (sticking to the substrate), indicating that the ZT3 coating was totally oxidized after 600 s ablation.  3.3. SEM observation  Representative surface and cross-sectional morphologies of all the coatings after 300 s ablation are shown in Fig. 7. The surface of the ZS3300 coating was composed of melted and solidified SiO2 and solid ZrO2 (Fig. 7a). SiO2 was also observed from the high magnification of the surface of the ZS2T1-300 and ZS1T2-300 coating (Fig. 7c and Fig. 7e). Besides, with the increasing of the TiC content, more melt area could be seen, especially for the ZT3-300 coating (Fig. 7g). The observation of the cross-sections of the ablated coatings indicated that a layered structure was formed after 300 s ablation for all the specimens. Combined with the structure and elemental mapping, different layers could be distinguished. Similarly, the bottom layers of all the ablated coatings were the un-oxidized layer for little O was detected. Therefore, the oxide scales were focused on. For the ZS3-300 coating, layer 1 was composed of glassy SiO2 and ZrO2 skeleton, presenting porous structure. And Layer 2 was supposed to be the SiC-depleted layer, which were confirmed by the Si mapping results (Fig. 7b). For the ZS2T1-300 coating, two layers were also observed for the oxide scale: layer 1 was similar to that of the ZS3-300 coating, which contained many pores. Layer 2 owned a relatively dense structure and almost no SiC-depleted area were detected. The oxide scale of the ZS1T2-300 coating was only composed of one layer, which seemed relatively dense (Fig. 7f). The microstructure of the ZT3-300 coating was similar to that of the ZS1T2300 coating. The cracks in Fig. 7(f-h) could be ascribed to the thermal  Fig. 3. Surface (a-d) and cross-sectional (e-h) morphologies of as-sprayed ZS3, ZS2T1, ZS1T2 and ZT3 coatings.  3  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 4. Macrographs of ZS3, ZS2T1, ZS1T2 and ZT3 coatings after ablation tests.  EDS results of the ZS3-600 coating. Almost no glassy SiO2 were observed which was confirmed by the EDS analysis (Fig. 8a). From the cross-sectional morphologies and elemental mappings, three distinct layers could be seen. Layer 1 and layer 2 belonged to the oxide scale and layer 3 was the un-oxidized scale. There was an crack in layer 2, which may be caused by the combination of the active oxidation of SiC and thermal stresses from the cooling process. The fractured morphologies of the ZS3-600 coating indicated that the structure of layer 1 was relatively porous and layer 2 was relative dense (Fig. 8(c-f)). Fig. 9 shows the surface, cross-section and related EDS results of the ZS2T1-600 coating. After 600 s ablation, more melted areas could be observed on the coating surface (Fig. 9a and b). The EDS results indicate that the melts were composed of Zr, Ti and O elements. From the cross-sectional morphology (Fig. 9c), it can be seen that the oxide scale could be divided into two layers. Layer 1 contained more pores compared with layer 2. The fractured morphologies of the ZS2T1-600 coating are given in Fig. 10. It can be seen that the area 1 of layer 1 was composed of particles associated with liquid phases (Fig. 10b). Layer 2 contained some threadlike pores (Fig. 10d) and layer 3 was the unoxidized layer (Fig. 10e). The surface, cross-section and related EDS results of the ZS1T2-600 coating are shown in Fig. 11. Almost no SiO2 could be seen on the surface (Fig. 11a). It is intersting to find that the oxide scale presented two distinct layers, which were different from that of the coating after  Fig. 5. Surface temperature curves of ZS3, ZS2T1, ZS1T2 and ZT3 coatings during ablation processes.  stresses during cooling process. The microstructure of the coatings after 600 s ablation were characterized in detail, trying to illustrate the effect of ablation time. Fig. 8 shows the surface, cross-sectional, fractured morphologies and related  Fig. 6. XRD patterns of ZS3, ZS2T1, ZS1T2 and ZT3 coatings after 300 s (a) and 600 s (b) ablation.  4  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 7. Surface, cross-sectional morphologies and related EDS results of ZS3-300 (a-b), ZS2T1-300 (c-d), ZS1T2-300 (e-f) and ZT3-300 (g-h) coatings.  300 s ablation. The difference of layer 1 and 2 might be resulted from the thermal gradient in the thick oxide scale. Moreover, there was a Ti rich area on the top surface, which exhibited a dense structure and might play as an effective oxygen diffusion barrier (Fig. 11c). The fractured morphologies of the ZS1T2-600 coating are illustrated in Fig. 12. It is worth mentioning that the crack in Fig.11c was caused by polishing because there was almost no crack can be observed from the fractured section (Fig. 12a). At area 1 in layer 1, particles associated with melts could be observed. Importantly, there was large amount of SiO2 retained in area 2, which was confirmed by the EDS result. Some  threadlike pores distributed in layer 2 (Fig. 12b) and layer 3 was the unoxidized coating (Fig. 12c). The thickness of the oxidized and un-oxidized layers of all the ablated coatings was measured and summarized in Table 2. It can be seen that the ZS1T2 coating exhibited the best oxidation resistance among all the coatings, having impact macrostructure and the largest thickness of the un-oxidized coating after 600 s ablation. From the above observations, the following phenomena are interesting and should be noted: (i) The ternary ZrC-SiC-TiC (ZS2T1 and ZS1T2) coating exhibited excellent ablation resistance compared with the ZS3 and ZT3 coatings;  Fig. 8. Surface (a), cross-sectional (b),  fractured morphologies (c-f) and related EDS results of ZS3-600 coating.  5  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 9. Surface (a-b), cross-sectional (c) and related EDS results of ZS2T1-600 coating.  (ii) the SiC-depleted layer was absent in the ablated ternary coatings.  4. Discussion  The ablation process is a combination of chemical corrosion and mechanical scour, leading to the loss of mass, the vaporization of liquid products and the escape of gaseous products. Several factors should be clarified before illustrating the failure mechanisms of the composite coatings: firstly, the difference in compositions, which would lead to different oxidation products; secondly, what is more important, the features and changes of the oxidation products during the ablation processes, like thermal stability, viscosity, oxygen permeability and so on. These factors were emphatically discussed as following.  4.1. Oxidation behaviors of ZrC, SiC and TiC  The chemical reactions of ZrC, SiC and TiC were influenced by both temperature and local oxygen pressure. For example, the oxygen partial pressure of layer 3 was significantly lower than that of the above layers. The following oxidation reactions may take place during the ablation processes [18,28]:  ZrC(s) + 2O2(g) → ZrO2(s) + CO2(g)  ZrC(s) + 3/2O2(g) → ZrO2(s) + CO(g)  SiC(s) + 2O2(g) → SiO2(l) + CO2(g)  SiC(s) + 3/2O2(g) → SiO2(l) + CO(g)  SiC(s) + O2(g) → SiO(g) + CO(g)  TiC(s) + 3/2O2(g) → TiO2(l) + CO(g)  TiC(s) + 2O2(g) → TiO2(l) + CO2(g)  TiC(s) + O2(g) → TiO(g) + CO(g)  (2)  (3)  (4)  (5)  (6)  (7)  (8)  (9)  4.1.1. Thermal dynamical prediction of oxidation products The volatility diagrams (the vapor pressure of predominant gaseous species, as a function of the equilibrium partial pressure of oxygen) of ZrC, SiC and TiC were constructed to predict the possible oxidation products [29,30]. The volatility diagrams of ZrC, SiC and TiC at 2300 °C were constructed by Factsage 7.2 software based on NIST-JANAF database [31], as shown in Fig. 13, aiming to predict the oxidation behavior of the ZS1T2 coating. The pO2 values for the generation of different condense phases of ZrC, SiC and TiC were calculated and summarized in Table 3. It can be seen that when 10−5 < pO2 < 100 Pa, there was a competition between active oxidation of SiC and oxidation of TiC. The XPS results of the inter layer of the ZS1T2-600 coating were supplied, as shown in Fig. 14. The XPS survey spectrum exhibited the presence of C, Zr, Si, Ti, and O elements (Fig. 14a). The high-resolution XPS spectrum of Ti 2p was well convoluted into four peaks (Fig. 14b): the binding energy at 458.1 eV and 464.1 eV was closed to that of Ti3+ 2p3/2 and Ti3+ 2p1/2 (457.9 eV, 464.0 eV), respectively [32,33]. While the other two peaks located at 459.1 eV and 464.9 eV were ascribed to Ti4+ 2p3/2 and Ti4+ 2p1/2 (458.9 eV, 464.5 eV), respectively. The presence of Ti3+ confirmed that there was some amount of Ti2O3 in the interlayer of the ZS1T2-600 coating. Combined with the absence of the SiC-depleted layer and the presence of Ti2O3, it could be induced that the competition of oxidation of TiC and active oxidation of SiC occurred during the ablation process, and the existence of Ti2O3 suppressed the formation of the SiC-depleted layer. A dense Ti-rich layer was observed in the ZS1T2-600 coating (Fig. 11c). Similar phenomenon had been observed by D.B. Lee [34], where a dense TiO2 layer was found for Ti3SiC2 after oxidation for 48 h in air at 1100 °C. The generation of the Ti-rich layer may be ascribed to the following two reasons: (i) TiO2 was less likely to evaporate or decompose, then retained to formed (Zr, Ti)O2 with ZrO2. While SiO2 was almost completely consumed after long time ablation. (ii) On the other hand, silicon ions in oxides could be relatively immobile, because of its  Fig. 10. Fractured morphologies of ZS2T1-600 coating.  6  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 11. Surface (a-b), cross-sectional (c) and related EDS results of ZS1T2-600 coating.  Fig. 12. Fractured morphologies and related EDS results of ZS1T2-600 coating.  Table 2  Thickness of oxide scale and un-oxidized coating at different ablation period.  Ablation time (s)  ZS3 (μm)  ZS2T1 (μm)  ZS1T2 (μm)  ZT3 (μm)  As-sprayed 300 600  Oxide  -  165 ± 10 295 ± 10  Un-oxidized  310 ± 10 210 ± 10 90 ± 15  Oxide  -  180 ± 5 390 ± 5  Un-oxidized  360 ± 10 255 ± 10 110 ± 10  Oxide  -  320 ± 5 340 ± 5  Un-oxidized  Oxide  Un-oxidized  390 ± 10 280 ± 10 180 ± 10  -  330 ± 10 Peeled  420 ± 10 290 ± 10 0  Fig. 13. Volatility diagrams of SiC (a), TiC (b) and ZrC (c) systems at 2300 °C.  higher bonding energy of Si+4-O (465 kJ mol−1) as compared with Ti+4-O (323 kJ mol−1) [34]. Therefore, it is possible for Ti4+ to gather in the outside layer and generate the Ti-rich layer. The consumption of SiO2 may be retarded because of the dense Ti-rich layer.  4.2. The features of SiO2 and TiO2  oxide scale and improving the oxygen diffusion resistance. Thermodynamic stability, viscosity and oxygen permeability are considered to be key factors for these phases, which would influence the ablation resistance of the coating and the microstructure of the oxide scale. The vapor pressure and viscosity of SiO2 and TiO2 were calculated and discussed as following:  ZrO2 was the skeleton of the oxide scale, and melted SiO2, TiO2 and (Zr,Ti)O2 played a role in associating the ZrO2 particles, densifying the  4.2.1. Thermodynamic stability During the ablation process,  SiO2  and TiO2 may  be  consumed  7  \\x0c', 'X. Pan, et al.  Table 3  Thermodynamic predictions of ZrC, SiC and TiC.  Corrosion Science 170 (2020) 108645  pO2 (Pa)  Predicted condense phase  Predicted gaseous phase  Predominant reaction  < 10−7  10−710−6 10−610−5 10−510−3  10−310° 100101 > 101  ZrC ZrC ZrC ZrC ZrO2 ZrO2 ZrO2 ZrO2  SiC SiC Si - - -  SiO2 SiO2  TiC TiC TiC TiC TiO Ti2O3 Ti3O5 TiO2  ZrC CO CO CO CO CO,CO2 CO2 CO2  SiC CO, Si, SiO CO, Si, SiO CO, SiO CO, SiO CO, SiO CO2, SiO2 CO2, SiO2  TiC CO, Ti, TiO CO, Ti, TiO CO, Ti, TiO CO, Ti, TiO CO, TiO CO2, TiO2 CO2, TiO2  - -  6 3, 6, 9 2, 3, 6 2, 4 2, 4, 8  through reaction (10−13).  SiO2(l) → SiO2(g)  SiO2(l) → SiO(g) +1/2O2(g)  TiO2(l) → TiO2(g)  TiO2(l) → TiO (g) +1/2O2(g)  (10)  (11)  (12)  (13)  The vapor pressure and decomposition pressure of SiO2 and TiO2 was calculated, as shown in Fig. 15. The vapor pressure and decomposition pressure of SiO2 were higher than that of TiO2, which meant SiO2 would be consumed easier during ablation processes. Besides, the calculated vaporization rate of SiO2 (207 mm/s) was nearly 900 times higher than that of TiO2 (0.23 mm/s) at 2225 °C [35], indicating a fast consuming of SiO2 in ultrahigh temperature environments.  4.2.2. Viscosity and oxygen permeability The viscosity of liquid phase has important influence on the mechanical scour resistance of the oxide scale. It is well known that viscosity as a function of temperature deviates from the Arrhenius relation [36]:  ln (η)  A  +  =  E RT  (14)  where η is the viscosity (dPa·s), A is the pre‐exponential factor (dPa·s), E is the activation energy (J mol−1), R is the gas constant and T is the absolute temperature. The viscosity of SiO2 and TiO2 were calculated by Factsage 7.2 using Viscosity module, as shown in Fig. 16. It can be seen that the viscosity of SiO2 was much higher than that of TiO2. Si-O tetrahedron could connect with each other and form Si-O network, while the network of Ti-O could hardly form at very high temperature [37]. Diffusivity of oxygen in liquid is inversely proportional to the viscosity of the liquid according to Stokes-Einstein relationship:  D=kT/6πηr  (15)  where D is diffusion constant [m2·s−1], k is Boltzmann’s constant [J·K−1], T is absolute temperature [K], η is viscosity [m2·s−1] and r is spherical particle radius [m]. It can be concluded that the oxygen diffusion rate in SiO2 was much lower than that in TiO2.  4.3. Ablation mechanism analysis of  the coatings  Distinguished ablation behaviors were observed for the ZS3, ZS2T1, ZS1T2 and ZT3 coatings in the present ablation conditions. The failure mechanisms of the four kinds of coatings could be concluded as follow:  4.3.1. For ZS3 coating SiO2 is characterized of low thermo stability, presenting high vapor and fast vaporization rate, which was tended to evaporated in a fast speed above 2000 °C. Therefore, the SiO2 in the ZS3 coating was consumed. The porous outer oxide layer would provide plenty of channels for oxygen diffusion. Besides, the active oxidation of SiC inside the coating was inevitable, and the significant generation of CO (g) and SiO (g) could lead to the creation of cracks/pores inside the coating (Fig. 8b), which caused the peeling of the oxide layers. Therefore, the ablation resistance of the ZS3 coating was not so good at environments above 2000 °C and the life-time is limited.  4.3.2. For ternary composite coating Due to the joint addition of TiC and SiC, the competition between the active oxidation of SiC and formation of Ti-contained phases helped a lot in avoiding the formation of SiC-depleted layer for both the ZS2T1 and ZS1T2 coatings. Thus, the ternary composite coatings maintained integrated. The variation of the contents of SiC and TiC additions had influence on the ablation behaviors. The consumption of SiO2 leaving porous oxide scale and make the ZS2T1 coating suffer severe oxidation  Fig. 14. XPS spectra of ZS1T2-600 coating: (a) survey spectra and (b) Ti 2p.  8  \\x0c', 'X. Pan, et al.  Corrosion Science 170 (2020) 108645  Fig. 15. Evaporation pressures and decomposition pressures of SiO2 and TiO2 at different temperatures.  Data availability  All research data supporting this publication are directly available within this publication.  CRediT authorship contribution statement  Xiaohui Pan: Writing original draft, Writing review & editing. Yaran Niu: Project administration, Supervision, Writing review & editing. Xueting Xu: Resources, Data curation. Xin Zhong: Resources, Data curation. Minhao Shi: Methodology. Xuebin Zheng: Project administration, Writing review & editing, Supervision. Chuanxian Ding: Supervision.  Declaration of Competing Interest  Fig. 16. Viscosity of SiO2 and TiO2 at different temperatures.  during the long-time ablation process. While for the ZS1T2 coating, the generation of a Ti-rich layer on the coating surface reduced the consumption of SiO2, endowing this coating the best ablation resistance among all kinds of coatings.  4.3.3. For ZT3 coating The oxidation resistance of the ZT3 coating was the worst among all the composite coatings. Although the coating kept intact after 300 s ablation, due to the high oxygen permeability of ZrO2 and TiO2, the ZT3 coating was completely oxidized after 600 s oxidation.  5. Conclusions  ZrC-SiC-TiC ternary composite coatings were designed and fabricated, which exhibited much better long-time ablation resistance than the ZrC-SiC and ZrC-TiC coatings in environments above 2000 °C. The anti-ablation property of ZrC-10 vol.%SiC-20 vol.%TiC coating was the best under the present ablation condition. The competition between the oxidation of TiC and SiC partially avoided the generation of the SiCdepleted layer. The Ti-rich layer formed on the coating surface played an important role in resisting oxygen diffusion and SiO2 consumption, which contributed to the improved anti-ablation property of the ternary composite coating. This work convinced that the design of ternary composite coating would be a promising way to enhance the long-life ablation resistance of ZrC in the environment above 2000 °C.  9  they have no conflicts of  The authors declared that work. We declare that we do not have any commercial or associative interest that represents a conflict of interest in connection with the work submitted.  interest  to this  Acknowledgements  This work was supported by the National Natural Science Foundation (for Young Scholar) of China under Grant 51102267 and Youth Innovation Promotion Association CAS (2014223).  References  [3]  [2]  [1] R. Savino, L. Criscuolo, G.D. Di Martino, S. Mungiguerra, Aero-thermo-chemical characterization of ultra-high-temperature ceramics for aerospace applications, J. Eur. Ceram. Soc. 38 (2018) 2937-2953. S. Xuetao, L. Kezhi, L. Hejun, D. Hongying, C. Weifeng, L. 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},{
  "_id": 119,
  "PDF": "Long-term oxidation behavior and mechanical strength degradation of a pressurelessly sintered ZrB2–MoSi2 ceramic.pdf",
  "Text": "['Scripta Materialia 53 (2005) 1297-1302  www.actamat-journals.com  Long-term oxidation behavior and mechanical strength  degradation of a pressurelessly sintered ZrB2-MoSi2 ceramic  Diletta Sciti *, Myle`ne Brach, Alida Bellosi  CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy  Received 30 December 2004; received in revised form 1 April 2005; accepted 25 July 2005  Available online 8 September 2005  Abstract  Oxidation tests were carried out on a pressurelessly sintered ZrB2\\x0020 vol.% MoSi2 ceramic from 700 to 1400 °C up to 100 h in air. The effects of oxidation on microstructure and flexural strength are discussed.  Ó 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: Zirconium diboride; Particulate composite; Oxidation; Microstructure; Mechanical properties  1. Introduction  The microstructure  and mechanical  properties  of  ZrB2-based composites have been recently investigated [1-6] and these materials were found to be viable  candidates  for a variety of high-temperature structural  applications. The major  problems  encountered with  hot pressing at temperatures P1900 °C, whilst  the final  density of pressurelessly sintered ZrB2-SiC typically does not exceed 75-80%. On the basis of these considerations,  the identification of a secondary phase that provides full  densification without  the need of pressure and that  im proves the oxidation resistance of ZrB2, would be a great advance from a technological point of view.  densification  [1-5]  and  One candidate material that may fulfill these require ZrB2-based materials high-temperature oxidation  involve  [7].  Sintering  powders  to  full  density  is  difficult,  unless  assisted sintering procedures 2000 °C  than  used  are  (if  and temperatures higher  no  sintering  additives  are used), due to the refractoriness of ZrB2 (Tmelting \\x183200 °C). The addition of sintering aids, like Ni [1], Si3N4 [2] or AlN [3], improves sinterability, but the formation of secondary phases can affect the oxida of  ZrB2 pressure ments is MoSi2. As a result of its thermodynamic instability, MoSi2 forms a surface layer of silica [11], which can help the densification and act as a protective barrier  against  oxidation  high-temperature  ZrB2. Dense materials were in fact produced by pressureless sintering at 1830 °C with an amount of 20 vol.% MoSi2 [12]. Furthermore, this system was found to withstand temperatures as high as 1400 °C during 30 h isothermal runs in  of  tion resistance of  the final material. The  addition of  synthetic  air without  significant  oxidation  [12]. Hot silicon carbide is highly beneficial for improving ZrB2 temperatures >1200 °C, due to  oxidation resistance at  the formation of a protective borosilicate glass  [8-10].  However, fully dense materials can be obtained only by  pressed ZrB2-MoSi2 ceramics have been studied and it was found that the addition of 10-30 vol.% MoSi2 improved both mechanical properties [13] and oxida tion  resistance  [13,14]  compared  to monolithic ZrB2  materials.  * Corresponding  author. Tel.: +39  0546  699748;  fax: +39  0546  46381.  E-mail address: dile@istec.cnr.it (D. Sciti).  In  this  contribution,  dense ZrB2 20 vol.% of MoSi2 particulate-composite was obtained by pressureless sintering. The oxidation resistance was  containing  fully  1359-6462/$ see front matter Ó 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2005.07.026  \\x0c', 'studied in the  temperature  range  700-1400 °C, up to  100 h in a  conventional  furnace. The microstructural  evolution was analyzed and discussed and the degrada tion of strength after oxidation was tested by four-point  bending.  2. Experimental procedure  Commercial powders were used to prepare ceramics  with  a  composition  of  80 vol.% ZrB2 MoSi2: ZrB2 Grade B (H.C. Starck, Germany), grain size range 0.1-8 lm, oxygen content 1 wt.% and MoSi2 (Aldrich, Germany) with mean particle size 2.8 lm and  and  20 vol.%  oxygen content of about 1 wt.%.  The powder mixture was milled for 24 h in absolute  ethanol using zirconia balls, subsequently dried in a rotary evaporator and sieved to \\x0060 mesh screen size. Sample bars with green dimensions 4 · 5 · 30 mm were  prepared by uniaxial pressing followed by cold isostatic  pressing under 350 MPa. The bars were pressurelessly  sintered in a resistance heated graphite furnace under a flowing argon atmosphere at 1830 °C for 30 min and the final densities were measured by the ArchimedesÕ  method. The samples were then oxidized in a conven tional  furnace in batches of and 1400 °C for  five  at  700,  800,  1000,  1200,  1300  100 h  in laboratory  air.  The weight gain/unit area was evaluated on all samples  by weight measurements  before  and  after  the  each  oxidation test with an accuracy of 0.01 mg.  Crystalline  phases  present  on  as-sintered  and  oxidized sample surfaces were detected by X-ray diffrac tion  (Cu  Ka  radiation, Miniflex  Rigaku,  Japan).  Surfaces and polished cross-sections were analyzed by  scanning electron microscope (SEM, Cambridge S360,  UK) and energy dispersive microanalysis (EDS, Model  INCA energy 300; Oxford Instruments, UK).  Room temperature  flexural  strength  of  as-sintered  and oxidized samples was measured in four-point bend ing. The test fixture had a lower span of 20 mm and an  upper span of 10 mm. A cross-head speed of 0.05 mm/  min was used on a universal  screw-driven test  frame  (Instron,  1195 USA). Five bars were  tested for  each  oxidation temperature.  3. Results  3.1. Microstructure of  the as-sintered material  The material was almost fully dense (relative density  > 99.5%). According to X-ray diffraction analysis,  the  crystalline phases  in the material bulk were ZrB2 and MoSi2. The microstructure (Fig. 1) of the polished surfaces had globular ZrB2 grains, that exhibited different contrast due to their different crystallographic orienta tions. The MoSi2 low dihedral angles, which accounts  phase was  characterized  by  very  for a high wetta bility towards the matrix phase during sintering. Minor (\\x185 vol.%) Mo-B, Mo-Si-B and Zr-C-O phases with variable stoi amount  of  secondary  phases  such  as  chiometry were detected by EDS analysis (see spectra in  Fig. 1). On the cross-sections of the bars, agglomerated  SiC grains were observed, mainly concentrated in the  near-surface regions of the bars to a depth of about 100 lm, whilst XRD diffraction patterns  revealed the  presence of ZrC and MoB on the surface of as-sintered  samples. All these secondary phases are thought to orig inate from reactions between the species constituting the  starting powders in presence of carbon and oxygen [12].  The sintering environment was,  in fact, carbon rich, due  to graphite heating  elements,  crucibles  and supports,  while  oxygen  impurities,  detected  by EDS  analysis,  could derive from the starting powders.  3.2. Microstructure modification induced by oxidation  The total weight gain/unit area measured on the oxi dized samples is shown in Fig. 2. Weight gain was very  limited  after  oxidation  in  the  temperature  range  Fig. 1. Polished surface of pressurelessly sintered ZrB2 + 20 vol % MoSi2 in back-scattered imaging and EDS spectra of selected areas.  1298  D. Sciti et al.  / Scripta Materialia 53 (2005) 1297-1302  \\x0c', '700-1000 °C, but  strongly increased when sample bars  were treated at higher temperatures. Crystalline phases  detected  by X-ray  diffraction  analysis  on  as-sintered  and on oxidized surfaces are presented in Table 1. The  intensity  of ZrB2 and MoSi2 increasing temperature of oxidation due to the masking  peaks  decreased with  effect of the growing oxide. The microstructural changes  induced  by  oxidation  are  described  for  each  tested  temperature. 700 °C. According to X-ray diffraction,  the  surface  crystalline phases consisted of monoclinic zirconia and  traces of MoO2. The  surface  (Fig.  3(a)) was  covered  by a discontinuous glassy layer  (containing Si, O and  several  impurities) and by aggregated zirconia particles.  Microcracking  of  the  surface  glassy  phase  probably  originated during cooling of  the glass, but  it cannot be  excluded that some of the cracks were due to oxidation  of  either molybdenum disilicides or MoB to gaseous  MoO3. The oxide 10 lm. 800 °C. The specimen surface (not  thickness  (Fig.  3(b)) was  less  than  shown) was cov ered either by a silica-rich layer or by a cracked layer  that had newly formed zirconia nano-crystals embedded  in it. The cross-section presented features those of the sample oxidized at 700 °C. 1000 °C. Monoclinic zirconia was, by far,  similar  to  the domi nant phase in the oxidation scale (Table 1), but a series  of minor crystalline phases were also detected,  including  MoB, ZrSiO4 and MoO2. The surface and section micrographs (\\x1812 lm) consisting of ZrO2 particles dispersed in amorphous silica containing vari (Fig.  4(a-c))  showed a compact  scale  ous  impurities and/or  traces of boron. Glassy bubbles  were  also visible on the  sample  surface. Beneath the  oxide layer, there 70-80 lm thick  was  a  subsurface  layer  about  containing  zirconia, Mo-oxide,  silica  pockets  and  cracks. Agglomerated  SiC  grains  and  MoB phases,  easily  visible  in the  cross-section, were  already  present  in  the  as-sintered material,  prior  to  oxidation. 1200 °C. The only crystalline phases detected on the  surface were monoclinic zirconia and zircon. A thick (\\x18150 lm) was observed on the and its composition was variable, containing mainly Si  oxide  sample  surface  and O, but also B and traces of other  impurities. The  surface morphology (not shown) resembled that of 1000 °C. The  the  sample oxidized at  image of  the  cross section (Fig. 5(a)) revealed a strong degradation of  the  sample.  Beneath  the  surface  layer  of  zirconia  and  silica-rich phase,  the material was heavily cracked to a 500 lm. Deep  depth  of  about  cracks  formed  in  the  oxidized specimens, mainly concentrated at the corners,  Fig. 2. Weight gain/unit area for 100 h oxidation tests at different  temperatures.  Table 1  Crystalline phases in the as-sintered material (bulk) and in the oxidized  samples  As-sintered  700  800  1000  1200  1300  1400  ZrB2  s  w  w  vw  MoSi2  w  w  vw  vw  m-ZrO2  s  vs  vs  vs  vs  vs  ZrSiO4  vw  vs  vs  s  MoO2  Traces  Traces  MoB  w  s  s  s  s  m-ZrO2 is monoclinic zirconia.  Fig. 3.  (a) Surface and (b) cross-section of a sample oxidized at 700 °C.  D. Sciti et al.  / Scripta Materialia 53 (2005) 1297-1302  1299  \\x0c', '1300  D. Sciti et al.  / Scripta Materialia 53 (2005) 1297-1302  Fig. 4.  (a-b) Surface and (c) cross-section of a sample oxidized at 1000 °C.  Fig. 5.  (a) Cross-section of a sample oxidized at 1200 °C, in back-scattered imaging, (b) detail of micrograph (a) showing boron-containing glass and  (c) related EDS spectrum (the C peak is due to graphite coating).  where penetration of  liquid boria  and/or borosilicate  glass was observed (Fig.  5(b))  and EDS spectrum in  Fig. 5(c)). B2O3 was not detected by X-ray diffraction probably owing to its amorphous nature. The cracks  glass. The  one,  cross-sections  two layers. The outer  (Fig. 6(b) and (c)) showed (thickness \\x18100 lm) was the silica-based oxide, which was heavily cracked and contained large zirconia grains (10-15 lm) with colum are thought to originate from stress concentration ,that  nar morphology, and Mo-B grains (Fig. 6(c)). The inner  developed due to the thermal expansion mismatch bet layer  had  a  thickness  of  about  0.8-1 mm,  contained  ween the oxide  layer and the underlying ZrB2-MoSi2  material. 1300 °C and 1400 °C. The phenomena involved were the sample oxidized at 1200 °C, but  similar to those of  partially  oxidized ZrB2 infiltrated. In addition,  and  cracks, where  the  glass  large  cracks  formed  at  the  corners of  the bars, as already explained.  at  these  temperatures,  the  oxidation  attack  involved  3.3. Flexural strength  most  of  the  sample  volume. The  surface  (Fig.  6(a))  was completely covered with a network of ZrSiO4 and monoclinic zirconia crystals, embedded in the silicate  The flexural  for unoxidized ZrB2- MoSi2 and for samples oxidized at various temperatures  strength values  (a) Surface, (b) cross-section images of a sample oxidized at 1300 °C and (c) enlarged view of the external part of the oxide layer in the cross Fig. 6.  section.  \\x0c', 'are presented in Fig. 7. The strength of  the as-sintered  sample was 372 ± 44 MPa. The changes in strength of samples oxidized up to 1200 °C were small and all within  the standard deviation of the as-sintered sample strength. A severe decrease (\\x1850%) was observed in bars oxidized at 1300 and 1400 °C.  4. Discussion  In order  to understand the oxidation behavior of  ZrB2-MoSi2 ceramic, one must first consider the oxidation behavior of the starting phases. Zirconium diboride  is well-known to oxidize according to the reaction [7]: ZrB2 þ 2:5O2 ðgÞ ! ZrO2 þ B2O3 ðlÞ  ð1Þ  Boric acid has a low melting point  (450 °C) and high  vapor  pressure, which make  it  vaporize  at  relatively  low temperatures: B2O3 ðlÞ ! B2O3 ðgÞ  ð2Þ  Below 1100 °C,  the  presence  of B2O3 barrier  as  a  liquid  is  known  to  act  as  an  effective  against oxygen T > 1100 °C,  transport  (parabolic  kinetic), whilst  at  B2O3 protective action (paralinear kinetics)  starts  to  vaporize  notably,  thus  reducing  its  [7].  MoSi2 oxidation follows the reaction [15,16]: MoSi2 ðsÞ þ 3:5O2 ðgÞ ! MoO3 ðsÞ þ 2SiO2 ðsÞ  ð3Þ  MoO3 has a high vapor pressure, which makes it volatile between 500 and 800 °C and it melts at 801 °C [16]. At T > 1200 °C,  the silica acts as a barrier against oxida tion. At lower temperatures, the oxidation of monolithic  MoSi2 is dominated by the formation of MoO3, which is much more rapid than the formation of SiO2 [17,18]. As a result, high internal stresses can arise at the grain  boundaries, due to volume expansion caused by MoO3 formation and volatilization.  In the present ZrB2-MoSi2 material, during oxidation at 700-1000 °C, the production of silica was too  slow to protect  the  surface. Therefore,  the dominant  process was the oxidation of ZrB2 to zirconia and liquid boria according to reaction (1), concurrent with the for mation and volatilization of MoO3 owing to reaction (3). Further oxidation involved the MoB and ZrC  secondary phases  formed during sintering, which were  mainly concentrated near the sample surface. According  to  thermodynamic  calculations, MoB can  oxidize  as  follows [19]: 2MoB þ 3:5O2 ðgÞ ! B2O3 ðlÞ þ 2MoO2  ð4Þ  Zirconium carbide experiences catastrophic oxidation at  relatively  low temperatures  owing  to  oxide  growth,  which causes high stresses along grain boundaries [20].  The oxidation reaction commonly reported is: ZrC þ 2O2 ðgÞ ! ZrO2 þ CO2 ðgÞ  ð5Þ  During oxidation of ZrB2-MoSi2 temperature range 1200-1400 °C, the silica production  specimens  in the  owing to reaction (3) should be fast enough to protect  the material.  Instead,  under  the  external  oxide,  a  strongly oxidized subsurface  layer was observed. One  possible explanation is that, at the early stage of oxida tion, when the silica oxide was not completely formed,  reactions  (1)-(5) caused a very rapid degradation. The  presence  of  fast-oxidizing  species  such  as MoB and  ZrC enhanced  the  process. The  formation  of  cracks  favored the  fast oxygen transport  to internal  regions  and increased the  area  exposed to oxidation. Liquid  boria reacted with silica forming a borosilicate  glass,  as previously  found for ZrB2-SiC composites The borosilicate glass, owing to its low viscosity, easily  [6,8,9].  penetrated into the bulk through open channels  and  voids  left behind by the oxidation of ZrB2. However, it is likely that the amount of glass present or its effective  thickness was not sufficient to fully protect the increased  area exposed to oxidation.  This  study highlights  the  fact  that  the addition of  20 vol.% MoSi2 was not enough to protect the ZrB2based material for long-term tests at T > 1200 °C. The  main factors determining the degradation between 1200 and 1400 °C, were the rapid subsurface oxidation  accelerated by the formation of cracks and the presence  of  fast oxidizing species  such as ZrC and MoB in the  near-surface region of the as-sintered specimens.  The flexural strength of  the samples was tested after  each oxidation. After thermal treatment at various tem peratures, the original flaw population of the as-sintered  material  changed, as a consequence of  the oxidation.  Despite the fact that strengths after oxidation at temperatures up to 1200 °C were not statistically different,  the  small increase of strength mean value after oxidation at 700 °C (compared to the as-sintered bars) was proba bly due to partial healing of the original critical flaws by  Fig. 7. Change in flexural strength in the as-sintered sample and after oxidation tests at 700-1400 °C.  D. Sciti et al.  / Scripta Materialia 53 (2005) 1297-1302  1301  \\x0c', 'the oxide. Up to 1200 °C, oxidation did not  seem to  have a detrimental effect on the strength of this material,  which was quite surprising since microstructural analy ses  revealed  remarkable microstructural  degradation.  The drop in strength observed after 1400 °C was  tests at 1300 and  directly  related  to  the  strong  damage  caused by the oxidation attack.  5. Conclusions  A ZrB2-based material 20 vol.% MoSi2 was produced by pressureless sintering. Oxidation tests carried out between 700 and 1400 °C for 100 h showed  containing  that this material was heavily oxidized at >1200 °C.  temperatures  Flexural  strength tests, performed on sample bars  after oxidation, highlighted that the strength was unaffected up to 1200 °C (\\x18370 MPa). The main factor determining the degradation of strength at higher  temperatures, was the rapid subsurface oxidation accel erated  by  the  formation  of  cracks  and  the  presence  of  fast-oxidizing  secondary  species  formed  during  sintering.  Acknowledgements  The work  is  supported  by  the European  Project  Research  Training  Network  HPRN-CT-2000-00044  ‘‘Composite Corrosion’’. The  research contract of M.  Brach is funded by the same project. The authors wish  to thank C. Melandri and S. Guicciardi for the measure ment of mechanical properties.  References  [1] Monteverde  F,  Bellosi A, Guicciardi  S.  J  Eur Ceram Soc  2002;22:279.  [2] Monteverde  F,  Guicciardi  S,  Bellosi  A. Mater  Sci  Eng  2003;A346:310.  [3] Monteverde F, Bellosi A. Adv Eng Mater 2003;7:508.  [4] Zhang GJ, Deng ZY, Kondo N, Yang JF, Ohji T. J Am Ceram  Soc 2000;83:2330.  [5] Chamberlain AL, Fahrenhotlz WG, Hilmas GE. J Am Ceram Soc  2004;87:1770.  [6] Opeka MM, Talmy IG, Wuchina EJ, Zaykosky JA, Causey SJ.  J Eur Ceram Soc 1999;19:2405.  [7] Tripp WC, Graham HC. J Electrochem Soc 1971;118:1195.  [8] Tripp WC, Davis HH, Graham HC. Ceram Bull 1973;52:612.  [9] Monteverde F, Bellosi A. J Electrochem Soc 2003;150:B552.  [10] Opila EJ, Halbig MC. Ceramic Engineering and Science Proceed ings, vol. 22. Westerville, (OH): The American Ceramic Society;  2001. p. 221.  [11] Jang YL, Lavernia EJ. J Mater Sci 1994;29:2557.  [12] Sciti D, Brach M, Bellosi A. J Mater Res 2005;20:922.  [13] Chamberlain AL, Fahrenholtz WG, Hilmas GE. Advances  in  ceramic matrix composites, ceramic transactions, vol. 153. Wes terville, (OH): The American Ceramic Society; 2003. p.299.  [14] Shaffer PTB. Ceram Bull 1962;41:96.  [15] Chou TC, Nieh TG. J Mater Res 1993;8:214.  [16] Zhu YT, Stan M, Conzone SD, Butt DP.  J Am Ceram Soc  1999;82:2785.  [17] Zhu YT, Shu L, Butt DP. J Am Ceram Soc 2002;85:507.  [18] Natesan K, Deevi SC. Intermetallics 2000;8:1147.  [19] HSC Chemistry for Windows 5, Outokumpu Research Oy, Pori,  Finland.  [20] Berkovich-Mattuck JB. J Electrochem Soc 1967;114:1030.  1302  D. Sciti et al.  / Scripta Materialia 53 (2005) 1297-1302  \\x0c']"
},{
  "_id": 120,
  "PDF": "Mechanical properties and ablation behavior of ZrB2-SiC ceramics fabricated by spark plasma sintering.pdf",
  "Text": "['Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  Contents lists available at ScienceDirect  Int. Journal of Refractory Metals and Hard Materials  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / I J R M H M  Mechanical properties and ablation behavior of ZrB2-SiC ceramics fabricated by spark plasma sintering  Xiang Zhang ⁎, Rutie Liu, Xiang Xiong, Zhaoke Chen  State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 29 May 2014 Accepted 11 August 2014 Available online 16 August 2014  Keywords:  Zirconium diboride Mechanical properties Ablation resistance  ZrB2-SiC ceramics were prepared by spark plasma sintering at 1375 °C for 5 min under a uniaxial load of 25 MPa. Mechanical properties and ablation behavior of ZrB2-SiC ceramics were investigated. The results showed that the flexural strength of ZrB2-30 vol.% SiC ceramic was significantly higher than that of ZrB2-15 vol.% SiC ceramic due to the larger amount of SiC whereas the ablation resistance was reduced. The increase in flexural strength was attributed to the sintering of SiC particles which located in the boundary of ZrB2 and enhanced the binding force between ZrB2 particles. The decrease in the ablation resistance at 3000 °C was due to the increase in weight loss and voids which were mainly caused by the rapid evaporation of SiO2. Ablation mechanism revealed that excessive SiC could supply much flaws and was hard to form continuous layer to seal the defects, which led to the dissatisfactory resistance to ablation for the ZrB2-30 vol.% SiC. The different ablation behaviors of central and outer ablation regions were also discussed.  © 2014 Elsevier Ltd. All rights reserved.  Introduction  Transition metal diborides including ZrB2, HfB2, TiB2 and TaB2 possess high melting temperatures in excess of 3000 °C, high thermal and electrical conductivity, as well as excellent chemical and physical stability at high temperatures [1-4]. Such properties make these diborides as the excellent candidates for thermal protection materials in both hypersonic vehicles and carbon/carbon composites in high temperature environment [5,6]. In addition, the existing available materials could not withstand extreme temperature of more than 2000 °C, which makes it necessary to develop the ultra-high temperature composites endowed with good ablation resistance [7]. Within the family of ultra-high temperature ceramics, ZrB2 has the lowest density and can form ZrO2 scale, which has attracted great attention in recent years [8]. However, the coexistence of covalent and metallic bond as well as low diffusivities becomes the barrier that impedes the process of sintering. Current researches have focused on densification, mechanical properties and oxidation behavior in the range of 1000 °C to 2600 °C, in which the ZrB2 ceramics were fabricated at a high temperature to overcome the poor sinterability [9-11]. In order to improve the mechanical properties and promote the sintering of ZrB2, sintering aids such as MoSi2, SiC, and ZrSi2 were added in ultra high temperature ceramics [12-14]. Numerous investigations indicate that SiC as the second phase in ZrB2 based ceramics not only enhances the mechanical properties, but also improves the oxidation resistance in the range of 1400 °C to 2000 °C through the obtainment of silica glass layer during  ⁎  Corresponding author. Tel.: + 86 15575186817; fax: + 86 731 88836079. E-mail address: zhangxiangcsu@csu.edu.cn (X. Zhang).  oxidation process [15-17]. The oxidation of UHTCs is a complex process, especially at temperature above 2000 °C. However, up to now, there are limited studies that reported the ablation mechanism and oxidation microstructure of ZrB2-SiC ceramics under oxyacetylene flame at about 3000 °C. The purpose of this paper is to report on the microstructure, density and mechanical properties as well as the ablation behaviors of ZrB2-SiC ceramics which were fabricated at 1375 °C. The effect of SiC on the ablation behavior was determined in terms of the ablation mechanism by investigating the weight loss and microstructure evolvement by means of the oxyacetylene torch at about 3000 °C for 100 s.  Experimental procedure  Commercially available ZrB2 powder (15 μm, N 99.5%, SBGT Optical Science and Technology Ltd., Beijing, China) and SiC (0.55 μm, N 99%, Shanghai Jing Chun Biochemical Technology Co., Ltd., Shanghai, China) were used as raw powders. The powder mixtures of ZrB2 plus 15 vol.% SiC (ZS15) and ZrB2 plus 30 vol.% SiC (ZS30) were ball mixed at 210 rpm for 20 h in stainless-steel ball media and ethanol, respectively. After mixing, the slurry was dried by vacuum oven and then screened by sieve of 80 mesh. The resulting powder mixtures were loaded into a 55-mm-diameter graphite die lined with graphite foil, and was spark plasma sintered (Model HPS-200, Kingtier New Alloy Material Co. Ltd., Anhui, China) in dynamic vacuum (b 1 Pa). The samples were heated from room temperature to 1375 °C at a rate of 100 °C/min and held for 5 min under a uniaxial pressure of 25 MPa. Finally, the samples were cooled down to room temperature in the furnace at a rate of 40 °C/min.  http://dx.doi.org/10.1016/j.ijrmhm.2014.08.006 0263-4368/© 2014 Elsevier Ltd. All rights reserved.  \\x0c', \"X. Zhang et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  121  The densities of the specimens were measured by the Archimedes method. Theoretical densities of the samples were calculated by applying the rule of mixture. Micro-hardness was measured by Vickers' indentation with a 0.98 N load applied for 15 s on polished sections. Flexural strength (σ) was tested in three points bending on 3 mm × 4 mm × 36 mm bars, using a 30 mm span and a crosshead speed of 0.5 mm/min. Each specimen was ground and polished with diamond slurries down to a 1 μm finish. The edges of all the specimens were chamfered to minimize the effect of stress concentration resulting from machining flaws. A minimum number of five specimens were tested in each experimental condition and the final result of mechanical properties of each ZrB2-SiC composite was the average of five samples. The microstructural observations and local element composition of the ZrB2-SiC ceramics were carried out by scanning electron microscopy (NOVA NanoSEM230, Czech Republic) along with energy dispersive spectroscopy (EDS, EDAX Inc.). The phase composition of the composite was determined by X-ray diffraction (Model D/max 2550V, Rigaku Ltd., Japan, Cu Kα radiation). ZS samples with a 21 mm × 21 mm × 5 mm cuboid for ablation testing were cut from the billet. The ablation test was carried out with an oxyacetylene flame system according to the standard of GJB 323-96A (China). The pressure and flux of O2 were 0.4 MPa and 0.42 L/s, and those of C2H2 were 0.095 MPa and 0.31 L/s, respectively. Coupons were placed in a direction vertical to the flame. The distance between the nozzle tip of the oxyacetylene gun and the central of the specimen was about 10 mm. The inner diameter of the nozzle tip was 2 mm. During the test, the ablation gun was ignited and the flame was steady, and the flame temperature was measured using an optical pyrometer and reached 3000 °C. Then the ablation gun was moved vertically to the surface of the specimen for 100 s. The mass ablation rates are obtained by a mass change of unit sample in unit time and are an average of 3 ablated samples. The mass ablation rate can be calculated by the following formula [18]:  Rm ¼ Δm=ðS \\x01 tÞ  ð1Þ  where Rm is the mass ablation rate, mg·cm-2·s -1; Δm is the mass change of the sample, mg; S is the surface area of the sample; t is the ablation time, s.  Results and discussion  Microstructure of the as-sintered ZS ceramics  Backscattered electron images of the polished surface of ZS15 and ZS30 are presented in Fig. 1. A white ZrB2 matrix with black dispersed SiC grains can be differentiated in ZS15 and ZS30. It was observed that the grain size of the SiC particles was about 3 μm to 5 μm in Fig. 1(a), and increased to about 10 μm in Fig. 1(b). Compared with the size of  Table 1 Mechanical properties of ZrB2-SiC composite.  Composition (vol.%)  Relative density (%)  Hv0.1 (GPa)  Flexural strength (MPa)  ZrB2-15 vol.% SiC ZrB2-30 vol.% SiC  93.1 94.2  10.5 11.1  335.5 391.6  SiC starting-powder (0.55 μm), a grain growth of SiC appeared at 1375 °C, which implied that the method of SPS could be conducive to sinter SiC. The SiC particles were primarily located at ZrB2 grain junctions and tended to be isolated. In contrast, the grain sizes of ZrB2 particles are about 10 μm to 15 μm, which does not have obvious grain growth during sintering.  Mechanical properties  The relative density and mechanical properties of samples ZS15 and ZS30 are listed in Table 1. It is found that the samples composed of two contents of SiC did not reach full density and the relative density of sample ZS15 is close to that of sample ZS30, which suggests that higher content of SiC could not obviously promote the densification of ZrB2 at this temperature. As seen from the table, the relative density, hardness and flexure strength of ZS30 increased as the SiC volume fraction increasing from 15 vol.% to 30 vol.%. The hardness of sample ZS15 (10.5 GPa) does not vary significantly comparing with that of ZS30 (11.1 GPa). The flexural strength of sample ZS30 (391.6 MPa) increased about 16.7% higher than that of sample ZS15 (335.5 MPa). It can be found that an increase content of SiC leads to a significant improvement of bending strength. The main reason might be the sintering processing, in which SiC particles located in the boundary of ZrB2 and enhanced the binding force between ZrB2 particles. The micrographs of fracture surface of samples ZS15 and ZS30 are shown in Fig. 2(a) and (b), respectively. Both of interand trans-granular fractures exhibited in samples ZS15 and ZS30.  Ablation properties  The ablation properties of ZrB2-SiC composites, including quality before ablation, mass loss and mass ablation rate are shown in Table 2. The results show that the mass ablation rate of ZS15 and ZS30 after 100 s ablation is 0.115 mg/cm2·s and 0.261 mg/cm2·s, respectively. Although it is demonstrated that ZS ceramic composite possesses good ablation resistance at ultra high temperature, the mass ablation rate of ZS ceramic composite was increased significantly as the content of SiC increased from 15 vol.% to 30 vol.%. The X-ray patterns of ZS samples after ablation are shown in Fig. 3. The result showed that the ablation product on the surface of ZS15 a n d Z S 3 0 w a s b o t h b a d d e l e y i t e , c o r r e s p o n d i n g t o J C PD S c a r d n o . 3 7 1 4 8 4 . It is known that baddeleyite contains 97%-98% ZrO2  Fig. 1. Typical backscattered electron images from a polished surface of ZrB2-SiC composites: (a) ZS15 and (b) ZS30.  \\x0c\", '122  X. Zhang et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  Fig. 2. SEM micrograph of the fracture surface of ZrB2-SiC composites: (a) ZS15 and (b) ZS30.  Table 2 Ablation properties of ZrB2-SiC composite.  Composition (vol.%)  Quality before ablation (g)  Mass loss (g)  Mass ablation rate (mg/cm2·s)  ZrB2-15 vol.% SiC ZrB2-30 vol.% SiC  9.4460 9.6339  0.0508 0.1151  0.115 0.261  and a little iron, in which iron is an impurity phase and introduced by ball milling with iron media, which might lower the sintering temperature of ZrB2 and resulted in a grayish yellow oxidation layer. Moreover, the other oxidative product such as SiO2 and B2O3 could not be identified because of the high evaporation under this extreme ablation condition. Fig. 4 shows the macroscopic photographs of ZS ceramic composite after 100 s ablation. The two distinct ablation regions (central ablation region and outer ablation region) can be found on the surface of the ablated samples. In the case of ceramic ZS15, the surface appeared to be a continuous ablation layer with a speckle pattern in the outer region and a splashing spot in the central ablation region. Concerning the surface of  ceramic ZS30, as shown in Fig. 4(b), a smooth ablation layer could be found in the outer ablation region and an evident crater in the central ablation region. After ablation text, the crater was surrounded by petaliform droplets, it could infer that the ZrO2 in the central ablation region melted and blew to be the droplets around the crater. While the peak of SiC and SiO2 could not be found in X-ray pattern (Fig. 3), the result showed that the oxidation of ZS30 had lower hightemperature stability than that of ZS15, which could be blow away and left a crater. Fig. 4(b) also shows that a crack was observed at the edge of ZS30. The formation of the crack was attributed to the internal stress, which was the formation of ZrO2 during ablation and the crystal transfer of ZrO2 during the period of cooling. Fig. 5 shows the surface microstructure and EDS results of ZS15 after ablation for 100 s. Some holes and droplets appeared on the surface due to the high air flux, high gas pressure and high evaporation rate of oxidative products beneath the outermost layer. The surface microstructure of outer ablation region and central ablation region of ZS15 sample was showed in Fig. 5(b) and (c), respectively. The outer ablation region contains a rough surface with holes, while the central ablation region owns a flat surface with cracks and holes. It should be noted  Baddeleyite(ZrO  )  2  (b)  (a)  15  20  25  30  35  40  45  50  55  60  65  70  75  80  2theta (degree)  Fig. 3. X-ray pattern of the two ZrB2-SiC composites: (a) ZS15 and (b) ZS30 after ablation for 100 s.  \\x0c', 'X. Zhang et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  123  Fig. 4. Surface appearance photos of ZrB2-SiC composites: (a) ZS15 and (b) ZS30 after 100 s ablation.  that the oxidation layer of ZrO2 in central and outer ablation regions acts as an armor to protect ceramic inside from flame scour and ablation. In accordance with XRD analysis, examination by EDS indicated that the surface of ablation center was composed of ZrO2. Thus, the cracks on the surface of central ablation region were resulted from the contraction stress from ZrO2, especially at the edge of the hole where suffered from the non-uniform stress. In addition, during the cool ing period, the crystal transition of ZrO2 would cause shearing and tensile stress in the oxidative surface and accelerate the formation of cracks. In the case of ZS30 under the same ablation test as shown in Fig. 6, evident difference could be observed from the surface compared with ZS15. Similar holes appeared in the same region but the sizes and  amounts of the holes in ZS30 increased significantly. As can be seen in Fig. 6(a), most of the holes were located in the central ablation region, which meant the central of sample underwent the severest oxidation. Fig. 6(b) and (c) shows the surface microstructure of the outer ablation region and central ablation region of ZS30 sample, respectively. The constituent in the outer ablation region of ZS30 was consumed a lot as well as the surface was transformed into skeleton structure. In contrast, a dense layer with crooked cracks and a ridge was formed in the central ablation region, which indicated that the ZrO2 in central ablation region melted and fused together. With regard to the distribution of element on the surface of ZS30, it could be found that the backscattering electron image and the EDS  Element  Wt%  OK  SiK  ZrL  Matrix  17.18  00.31  82.51  At%  53.98  00.55  45.47  Correction  ZAF  Fig. 5. SEM and EDS results of the surface of ZS15 after ablation at 3000 °C for 100 s. (a) High magnification image of part 1 in Fig. 4, (b) high magnification of outer ablation region, and (c) high magnification of central ablation region and EDS analysis.  \\x0c', '124  X. Zhang et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  Element  Wt%  OK  SiK  ZrL  Matrix  27.54  11.43  61.03  Correction  ZAF  At%  61.53  14.55  23.91  Element  Wt%  OK  SiK  ZrL  Matrix  14.25  00.43  85.32  At%  48.38  00.82  50.80  Correction  ZAF  Fig. 6. Surface micrographs and EDS results of ZS30 after ablation at 3000 °C for 100 s. (a) High magnification image of part 2 in Fig. 4, (b) high magnification of outer ablation region, (c) high magnification of central ablation region, and (d) backscattered electron morphology and EDS analysis of central ablation region.  ana lys is resu lts revea led the centra l ab la t ion region composed o f elements of Zr, Si and O. Fig. 6(d) has showed the distribution of different elements, and it was not hard to find the element of Si mainly existed in the gray area surrounded around the holes in central ablation region, which means that these holes allowed the oxygen to flow into the inner, and the oxidation of SiC such as the glassy and gaseous SiO2 would run out of the sample through the holes. Thus, the SiO2 could be remained around the holes due to the effect of surface tension of liquid SiO2 during the ablation test.  Effect of SiC on the ablation resistance  It is well known that the ablation under ultra high temperature is an extremely complex process including various oxidation reactions, the  erosion of high speed a ir scour and the evaporat ion of ox idat ive products. ZS ceramic composite exhibited a good ablation resistance under the extreme ablation condition at about 3000 °C, especially for the samples with the content of 15 vol.% SiC . According to the mass loss of ZS15 and ZS30 , a larger content of S iC leads to more mass loss. It could be inferred that the holes were formed in the location where it was filled with SiC before ablation. And the increase in content of SiC produced more voids on the surface of ablated samples which led a worse ablation resistant. In addition to the influence of higher content of SiC, analysis about the EDS results and microstructure of ZS30 as shown in Fig. 6(d) indicated that the oxidation of SiC was remained and formed into a new phase on the surface. Dong Gao et al. had studied the formation of zircon in the oxidation process of ZrB2-SiC composite [19]. Reported work has also  \\x0c', 'X. Zhang et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 120-125  125  noted that a dense and continuous layer of zircon could significantly improve the oxidation resistance performance of ZrB2-SiC [19,20] . Compared with Fig . 6(d) of ZS30 , it illustrated that the ZrO2-SiO2 could not form a continuous and dense layer to seal the defects such as cracks and holes under extremely high temperature (3000 °C). During the ablation of ZrB2-SiC composite under oxyacetylene flame, the following oxidation reactions may happen [2,16,19,20].  SiCðsÞ þ 2O2  ðgÞ→SiO2  ðgÞ þ CO2  ðgÞ  SiC sð  Þ þ 3  2  O2 gð Þ→SiO2 gð  Þ þ CO gð Þ  ZrB2  sð Þ þ 5  2  O2 gð Þ→ZrO2  lð Þ þ B2O3 gð  Þ  ð1Þ þ SiO2  ðgÞ→ZrSiO4  ð1Þ  ZrO2  ð1Þ  ð2Þ  ð3Þ  ð4Þ  The main chemical process during ablation for ZS composites is the ablatant obtained according to the oxidation reactions (Eqs. (1)-(3)), which produce an ablation layer that mainly consists of ZrO2. When the temperature was above the boiling point of SiO2 (about 2300 °C), the evaporation of SiO2 was serious and it was hard for SiO2 to remain on the surface of ZS [5,16]. And at this extremely high temperature, borosilicate combined by B2O3 and SiO2 would be difficult to form due to the rapid evaporation [2]. In contrast, the glassy and gaseous SiO2 as well as molten ZrO2 made it possible to form zircon (ZrO2-SiO2), which had a higher melting point and viscosity than borosilicate [16]. But the zircon (ZrO2-SiO2) is hard to form a continuous layer to seal the defects on the surface of ZS composite. The main reason according to the formation process of zircon reported by Dong Gao et al. might be the insufficient SiO and silica [19]. On one hand, the amount of oxidation of SiC was controlled by the temperature of oxyacetylene flame. Higher temperature will accelerate the oxidation rate. On the other hand, the flame scour and evaporation of silica will increase as the temperature rise up. So, it can be seen that the balance between evaporation and oxidation to remain enough oxidation of SiC is difficult to achieve. According to the former analyzes, it indicated that ZrB2-SiC composites could endure the erosion when high temperature reaches 3000 °C. The mass loss of ZS composites is mainly attributed to the oxidation , evaporation and subl imation of SiC which was blow off apace in the process of oxidation. With the increase in SiC content, such assistance to the growth and oxygen transport voids became more evidently due to the degraded structure. As concluded, the addition of SiC was advantageous to the mechanical properties of ZS composite, however, excessive SiC was evidently adverse to the ablation resistance at extremely high temperature. Consequently in perspectives combined with mechanical properties and ablation resistance, the addition of SiC to ZrB2 ceramic should be controlled appropriately.  Conclusions  ZrB2-15 vol.% SiC and ZrB2-30 vol.% SiC ceramic composites were fabricated by spark plasma sintering under a uniaxial load of 25 MPa  at 1375 °C for 5 min. The results showed that mechanical properties of the as-sintered ZS ceramics are mainly attributed to the dense and content of SiC. By means of the oxyacetylene torch, ablation resistance of ZS composites was investigated in terms of the effect of SiC. Results indicated that ZS15 provided the lighter oxidized structure and mass loss than ZS30. Particularly, the compact and continuous oxide scale with some holes in central and outer ablation region acts as an armor to protect ceramic inside from flame scour, which revealed that the ZS15 composite can be applied at about 3000 °C in oxide environment for more than 100 s.  Acknowledgments  This work was supported by the National Basic Research Program of China (2011CB605805) and the Fundamental Research Funds for the Central Universities of Central South University (2014zzts178).  References  [5]  [7]  [1] Ren X, Li H, Fu Q, Chu Y, Li K. TaB2-SiC-Si multiphase oxidation protective coating for SiC-coated carbon/carbon composites. J Eur Ceram Soc 2013;33:2953-9. [2] Zhou S, Li W, Hu P, Hong C, Weng L. Ablation behavior of ZrB2-SiC-ZrO2 ceramic composites by means of the oxyacetylene torch. Corros Sci 2009;51:2071-9. [3] Sciti D, Balbo A, Bellosi A. Oxidation behaviour of a pressureless sintered HfB2-MoSi2 composite. J Eur Ceram Soc 2009;29:1809-15. [4] Silvestroni L, Guicciardi S, Melandri C, Sciti D. TaB2-based ceramics: microstructure, mechanical properties and oxidation resistance. J Eur Ceram Soc 2012;32:97-105. Li H-j, Yao X-y, Zhang Y-l, Li K-z, Guo L-j, Liu L. Effect of heat flux on ablation behaviour and mechanism of C/C-ZrB2-SiC composite under oxyacetylene torch flame. Corros Sci 2013;74:265-70. [6] Sciti D, Savino R, Silvestroni L. Aerothermal behaviour of a SiC fibre-reinforced ZrB2 sharp component in supersonic regime. J Eur Ceram Soc 2012;32:1837-45. Li G, Han W, Zhang X, Han J, Meng S. Ablation resistance of ZrB2-SiC-AlN ceramic composites. J Alloys Compd 2009;479:299-302. [8] Zhang X, Hu P, Han J, Meng S. Ablation behavior of ZrB2-SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions. Compos Sci Technol 2008;68:1718-26. [9] Patel M, Singh V, Reddy JJ, Bhanu Prasad VV, Jayaram V. Densification mechanisms during hot pressing of ZrB2-20 vol.% SiC composite. Scr Mater 2013;69:370-3. [10] Han J, Hu P, Zhang X, Meng S, Han W. Oxidation-resistant ZrB2-SiC composites at 2200 °C. Compos Sci Technol 2008;68:799-806. [11] Guo W-M, Vleugels J, Zhang G-J, Wang P-L, Van der Biest O. Effect of heating rate on densification, microstructure and strength of spark plasma sintered ZrB2-based ceramics. Scr Mater 2010;62:802-5. [12] Guo S-Q. Densification of ZrB2-based composites and their mechanical and physical properties: a review. J Eur Ceram Soc 2009;29:995-1011. Ikegami M, Guo S, Kagawa Y. Densification behavior and microstructure of spark plasma sintered ZrB2-based composites with SiC particles. Ceram Int 2012;38: 769-74. [14] Guo S-Q, Kagawa Y, Nishimura T, Tanaka H. Pressureless sintering and physical properties of ZrB2-based composites with ZrSi2 additive. Scr Mater 2008;58:579-82. Liu Q, Han W, Zhang X, Wang S, Han J. Microstructure and mechanical properties of ZrB2-SiC composites. Mater Lett 2009;63:1323-5. [16] Hu P, Guolin W, Wang Z. Oxidation mechanism and resistance of ZrB2-SiC composites. Corros Sci 2009;51:2724-32. [17] Sarin P, Driemeyer PE, Haggerty RP, Kim DK, Bell JL, Apostolov ZD, et al. In situ studies of oxidation of ZrB2 and ZrB2-SiC composites at high temperatures. J Eur Ceram Soc 2010;30:2375-86. [18] Wang Y-l, Xiong X, Li G-d, Liu H-f, Chen Z-k, Sun W, et al. Ablation behavior of HfC protective coatings for carbon/carbon composites in an oxyacetylene combustion flame. Corros Sci 2012;65:549-55. [19] Gao D, Zhang Y, Fu J, Xu C, Song Y, Shi X. Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: formation of zircon. Corros Sci 2010;52:3297-303. [20] Yao X-Y, Li H-J, Zhang Y-L, Ren J-J, Yao D-J, Tao J. A SiC/ZrB2-SiC/SiC oxidation resistance multilayer coating for carbon/carbon composites. Corros Sci 2012;57:148-53.  [13]  [15]  \\x0c']"
},{
  "_id": 121,
  "PDF": "Mechanical properties and oxidation behavior of spark plasma sintered (Zr,Ti)B2 ceramics with graphene nanoplatelets.pdf",
  "Text": "['Ceramics International 46 (2020) 26109-26120  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Mechanical properties and oxidation behavior of spark plasma sintered (Zr, Ti)B2 ceramics with graphene nanoplatelets  T  Melis Kaplan Akarsu,  Ipek Akin∗  Istanbul Technical University, Department of Metallurgical and Materials Engineering, Maslak, 34469,  Istanbul, Turkey  A R T I C L E  I N F O  A B S T R A C T  Keywords: Borides Solid solution Spark plasma sintering Oxidation  This research focuses on the production and characterization of (Zr,Ti)B2 solid solution systems. The purpose of this study is to understand how solid solution formation influences the structure, mechanical properties, and oxidation resistance of ZrB2 and TiB2. For this, herein, monolithic ZrB2, TiB2, and (ZrxTi1-x)B2 solid solution systems (x = 0.25, 0.50, 0.58, 0.75) were prepared. Moreover, 1 vol% graphene nanoplatelets (GNPs) were added to two different solid solution systems (Zr0.25Ti0.75B2 and Zr0.58Ti0.42B2) without affecting the complete solid solubility of these systems. All powder mixtures were densified by spark plasma sintering at 1800 °C, under 50 MPa with 5 min holding time. The densification, oxidation behavior, microstructure, phases, and mechanical properties (Vickers hardness and indentation fracture toughness) of these samples were systematically investigated. The formation of solid solutions caused an increase in the Vickers hardness and fracture toughness of the samples. Further improvement in mechanical properties and better oxidation resistance were achieved with the addition of GNPs to the solid solution systems.  1.  Introduction  Elements, such as Ti, Zr, Nb, and Ta, of the fourth and fifth groups of the periodic table form the most stable borides [1]. Transition metal borides of the groups IVB and VB are members of a family of ceramic materials known as ultra-high-temperature ceramics (UHTCs) [2]. High melting temperatures, excellent chemical stability, high electrical and thermal conductivities, and good corrosion resistance make these ultrahigh-temperature borides potential candidates for use in extreme chemical and thermal environments [3]. These diborides are used in various potential applications such as in thermal protective structures for the leading-edge components of hypersonic re-entry space vehicles, propulsion systems, furnace elements, refractory crucibles, and plasma arc electrodes [4]. As the materials are exposed to high temperatures, high temperature stability and oxidation behavior are important properties for the application of diborides as engineering ceramics [5]. To date, various borides such as titanium boride (TiB2), zirconium boride (ZrB2), and hafnium boride (HfB2) have been investigated. As borides are exposed to high temperatures under working conditions, their use in monolithic forms is not suitable for achieving sufficient structural properties. Low fracture toughness and poor oxidation resistance also limit the use of monolithic borides. Therefore, the addition of Ni, Nb, Fe, V, C, disilicide structures (MoSi2 and TiSi2), B4C, Si3N4,  AlN, and SiC into monolithic borides and strategies for improving the other properties, especially densification, of borides have been carried out [6-8]. However, it has been reported that these secondary phase additions adversely affect some properties of metallic borides. For example, some metallic-based additives facilitate liquid sintering by forming a liquid phase; however, due to excessive grain growth, the trapped pores remain in the structure and high densities cannot be achieved [6]. Moreover, metallic additives reduce the resistance of borides to high temperatures and corrosion [8,9]. It is possible to sinter the borides to high theoretical densities using ceramic-based additives; however, the secondary phases added to the structure may settle at the boride grain boundaries and adversely affect the high-temperature properties [6]. Nevertheless, there are limited exceptions, especially WC, in this regard. It has been reported that the high temperature strength (~680 MPa at 1600 °C) of the ZrB2-20 vol% SiC composite increases with the addition of 5 vol% WC due to the removal of oxide impurities from grain boundaries [10]. Various studies on double diboride systems have also been reported in the literature, which indicate that metallic or carbide secondary phase additions do not adversely affect the superior properties of borides [3,11,12]. Moreover, structures with better densification and superior mechanical properties have been obtained via the formation of boride solid solutions than that of double boride systems [11,12].  ∗ Corresponding author. E-mail address: akinipe@itu.edu.tr (I. Akin).  https://doi.org/10.1016/j.ceramint.2020.07.106 Received 24 February 2020; Received in revised form 30 June 2020; Accepted 10 July 2020  Available online 27 July 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', \"M.K. Akarsu and I. Akin  Table 1  Ceramics International 46 (2020) 26109-26120  Compositional data, relative density, average grain size, Vickers hardness, and indentation fracture toughness of the samples.  Sample name  Composition  Relative density (%)  Average grain size (μm)  Vickers hardness (GPa)  Indentation fracture toughness (MPa·m1/2)  TiB2 25Zr-75Ti 25Zr-75Ti-G 50Zr-50Ti 58Zr-42Ti 58Zr-42Ti-G 75Zr-25Ti ZrB2  Monolithic TiB2  Zr0.25Ti0.75B2  Zr0.25Ti0.75B2 + 1 vol% GNP  Zr0.5Ti0.5 B2 Zr0.58Ti0.42B2  Zr0.58Ti0.42B2+ 1 vol% GNP  Zr0.75Ti0.25 B2  Monolithic ZrB2  95.26 98.90 99.74 99.02 99.33 99.32 98.55 98.01  19.7 ± 5.75 12.8 ± 4.21 15.2 ± 5.78 9.5 ± 2.82 15.4 ± 5.09 27.1 ± 8.51 17.8 ± 4.29 57.2 ± 9.08  22.71 ± 0.98 26.33 ± 0.95 23.52 ± 0.96 24.39 ± 0.42 23.43 ± 0.36 20.58 ± 0.73 18.73 ± 0.33 13.16 ± 0.44  3.93 ± 0.36 4.93 ± 0.19 5.04 ± 0.23 2.97 ± 0.32 2.69 ± 0.28 3.03 ± 0.42 2.66 ± 0.17 3.08 ± 0.26  Mondal et al. [13] have produced a ZrB2-TiB2 system by adding 5-30 wt% of TiB2 to ZrB2 via mechanical activation-spark plasma sintering at 2100 °C for 15-min dwell time under 50 MPa uniaxial pressure and an argon atmosphere and reported continuous solid solution formation for all compositions. It has been reported that densification, hardness (~20.6 GPa), fracture toughness (~4.7 MPa m1/2), and abrasion properties were improved for all the samples except for the system containing 5% TiB2. Phase analysis during solid solution formation is very important to interpret the behavior of the atoms inthe peaks (2θ value), changes in peak widths volved in the shifting of due to variation in the composition, lattice parameter changes, and variable percentages. The lattice parameters measured by XRD should agree with the values calculated by the Vegard's law. The Vegard's law is an approximation in ideal solutions for the two components that have less than 5% difference in lattice parameters. For a continuous substitutional solid solution, the unit cell parameters should vary linearly with composition [14]. To improve both the functional and mechanical properties of various ceramic matrices, graphene in the form of graphene nanoplatelets (GNPs) or graphene oxide (GO) is used as a filler [15]. Graphene has sp2 carbon hexagonal networks, where adjacent carbon atoms are strongly covalently bonded to each other [16]. Graphene is an ideal candidate as an advanced filler material for ceramic composites due to its excellent mechanical, thermal, and electrical properties. The interfaces between ceramics and graphene are hardly known, and further studies are required to understand the interface adhesion fracture energy, its dependence on crack speed and temperature, and the effects of processing temperatures and environments on it [17]. Previously, as ultra-hightemperature ceramics, graphene-reinforced transition metal boridebased ceramic composites have been used in the aerospace industry as high-temperature barriers for nose caps in space shuttles and military ballistic equipment [18]. To the best of our knowledge, this is the first study on the addition of graphene to a (Zr,Ti)B2 solid solution and investigation of the effects of graphene addition on the mechanical properties and oxidation behavior of the (Zr,Ti)B2 solid solution. Owing to the difficulty in the densification of solid solutions, spark plasma sintering (SPS) has been adopted [19]. Inagaki et al. [20] have reported that the solid solution reaction is enhanced by SPS. Furthermore, to avoid the degradation of GNPs, the reduction of SPS temperature and time is critical. SPS is a favorable rapid sintering technique that minimizes any degradation and is useful for the investigation of the sintering behavior of ceramics with nanofillers [21]. This study was aimed at producing a (Zr,Ti)B2 solid solution with better densification, oxidation behavior, and higher mechanical performance than the monolithic forms of borides by SPS. GNPs were added to the solid solutions to enhance the density, mechanical properties, and oxidation behavior of these solid solutions for their use in the nose cap and leading-edge of the hypersonic and atmospheric reentry systems.  2. Materials and methods  2.1. Powder preparation  The starting materials used herein were commercially available ZrB2 (H.C. Starck, Grade B; the impurity (wt.%) information provided by the supplier is as follows: C: 0.1-0.2, O: 0.8-1.5 and N: 0.06-0.25), TiB2 (H.C. Starck, Grade D; the impurity (wt.%) information provided by the supplier is as follows: C: 0.2-0.5, O: 1.1 and N: 0.4-0.6) and GNP powders (Nanokomp, purity > 97%, thickness: 5-8 nm, and diameter: 5-10 nm). The average particle sizes of the starting ZrB2 and TiB2 powders were measured as ~3.1 and ~5.9 μm, respectively, using a laser particle size analyzer (Malvern Mastersizer 2000). The ZrB2 and TiB2 powders were ball-milled together for 24 h inside a polythene bottle containing SiC balls in ethanol. The GNPs were separately ball-milled with YSZ balls in ethanol for 24 h. Then, each powder was dispersed in ethanol for 5 min by an ultrasonic homogenizer (Bandelin Sonopuls HD 2200, operating at 50% amplitude with on and off cycles). After the dispersion of the ZrB2-TiB2 powders and GNPs, the resulting samples were mixed. Magnetic stirring was applied for 5 h for a good dispersion of the powders before evaporating the ethanol. The powder mixture was dried in a drying oven at 100 °C for 24 h and then pounded in an agate mortar to obtain a soft and nonagglomerated starting powder. Cylindrical graphite die with an inner diameter of 50 mm was used for SPS. The entire contact zone of the powder with the punches and die was covered with a graphite sheet to ensure better conductivity.  2.2. SPS of  the samples  SPS was conducted using an SPS apparatus (7.40 MK-VII, SPS Syntex Inc.). A pulsed direct current (12 ms/on, 2 ms/off) was applied during SPS under a vacuum atmosphere. The SPS temperature, time, uniaxial pressure, and heating rate for the samples were 1800 °C, 5 min, 50 MPa, and 100 °C/min, respectively. The experimental compositions are summarized in Table 1. During the experiments, the temperature was measured and controlled using an optical pyrometer (Chino, IR-AH), which was focused on a small hole on the wall of the graphite die. SPS was conducted in the temperature-controlled mode while monitoring the shrinkage behavior of the samples.  2.3. Characterization of samples  The density of the solid solutions was determined from the measured Archimedes bulk density, where distilled water was employed as the immersion medium. The rule of mixtures was used for calculating the theoretical density of the samples. The phase compositions of the samples were determined by XRD (MiniFlex, Rigaku Corp.) in the 2θ range of 20°-90° at the scanning rate of 2°/min with CuKα radiation. The morphology of the samples was characterized by scanning electron microscopy (SEM, JSM 7000 F, JEOL Ltd.). Before SEM, the sintered samples were polished to remove the graphite sheet, and the fracture surfaces of the samples were  26110  \\x0c\", 'M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  investigated. The samples were coated with gold by sputtering. Chemical etching was performed for 5 s using a solution containing 30 mL lactic acid (90%), 10 mL nitric acid (65%) and 10 mL hydrofluoric acid (40%) [22]. The average grain size of the samples was measured by the linear intercept method using the ImageJ software. The microhardness of the samples was measured using a microhardness tester (VHMOT, Leica Corp.) fitted with a Vickers indenter. The samples were ground and polished before conducting the measurements. Indentations were produced on the polished surfaces under a load of 9.8 N. The indentation fracture toughness of each sample was calculated using the Anstis equation [23] based on the half-length of the crack formed around the indentations under a load of 19.6 N. Thermogravimetric (TG) analysis was performed in air using a TG/DTA analyzer (PerkinElmer Diamond TG DTA). The samples were heated from 25 °C to 1100 °C at the heating rate of 5 °C/min. The oxidation resistance of the sintered samples was tested in stagnant air at temperatures of 1100 and 1200 °C for 180 min using a MoSi2 resistanceheated furnace (Nabertherm C42). At the oxidation temperature, the samples were placed into the furnace and after the dwell time samples were free cooled to room temperature inside the furnace. The specific mass change was calculated by dividing the total mass change by the surface area of the samples. Raman microscopy was conducted to characterize the GNPs using a Bruker SENTERRA dispersive Raman spectrometer equipped with an Olympus confocal microscope with a 20× lens and a laser device operating at 785 nm excitation wavelength. The signals were recorded by the TE-cooled CCD detector. Zr0.25Ti0.75B2 and Zr0.58Ti0.42B2 were selected to determine the effect of GNP addition on their properties because these samples achieved the combination of highest densification and best mechanical properties.  3. Result and discussion  3.1. Densification  The relative densities of the solid solutions of monolithic borides with and without GNPs are shown in Table 1, which were found to be ~95.3% and ~98.0% for monolithic sintered TiB2 and ZrB2, respectively. The solid solution formation enhanced the densification of monolithic borides, and the formed samples exhibited higher relative densities than the monolithic samples. All solid solution samples achieved the relative densities of more than ~98%. Density measurements indicated that the samples had negligible open porosity. The measured densities of the samples increased with an increase in the amount of ZrB2 in the composition due to the higher theoretical density of ZrB2 (6.09 g/cm3). Highest relative density (99.3%) was obtained for Zr0.58Ti0.42B2. Densification enhancement with the formation of a complete solid solution was also reported by Mondal et al. [13]. Oxygen impurities affect the densification of borides (MeB2). MeO2 and B2O3 in equimolar amounts are formed under ambient conditions, and then, B2O3 evaporates at elevated temperatures. The remaining porous MeO2 can react with additives, such as B4C, MoSi2, and C, and remove oxygen impurities from the surfaces of the powder particles. In this way, the densification of borides can be improved, which causes a decrease in grain coarsening and improvement in the oxidation resistance [24]. Borothermal, carbothermal, and B4C reduction are the main reaction paths used for the reduction of MeO2 [25-27]. For the reduction reactions, several studies have utilized different carbon sources such as carbon black [28], graphite nano-flakes [29], graphene [30] and carbon nanotubes [31]. Upon the addition of 1 vol% GNPs, the relative density of the samples (~99.3%) was retained. Additionally, GNPs enhanced the densification of Zr0.25Ti0.75B2. The displacement graphs of Zr0.25Ti0.75B2 and 25Zr-75Ti-G (Fig. S1) as a function of temperature were used to obtain information about the densification of these samples. The temperatures at which the shrinkage of Zr0.25Ti0.75B2 started and completed decreased from 1654 to 1593 °C and 1780 to 1694 °C, respectively, with  the addition of 1 vol% GNPs. This could be attributed to the carbothermal reduction of ZrO2 and TiO2 by the GNPs and removal of oxygen impurities from the structure, preventing the entrapment of pores and causing nearly full densification of 25Zr-75Ti-G and 58Zr-42Ti-G. The GNPs react with ZrO2 and TiO2 according to Eqs. (1) and (2), which become favorable above ~1500 °C in the standard state. Due to the vacuum atmosphere of SPS, Eqs. (1) and (2) became favorable at significantly lower temperatures, and CO(g) was removed from the chamber [30,32,33]. During the reaction, liquid B2O3 vaporizes at ~1450 °C according to Eq. (3) [34].  ZrO  2 ( s  )  +  TiO  2 ( s  )  +  B O  2  3 ( ) l  B O  2  3 ( ) l  5 C  +  →  ZrB  2 ( s  )  +  5 CO g (  )  5 C  +  →  TiB  2 ( s  )  +  5 CO g (  )  B O  2  3 ( ) l  B O  2  3 ( g  )  →  (1)  (2)  (3)  Moreover, it is known that GNPs initiate particle rearrangement in the early stage of sintering due to their self-lubricating effect and act as a sintering aid in the later stages due to their carbon structure [35]. Additionally, GNPs can enhance the electrical and thermal conductivity of powders. The enhanced conductivity provides more uniform heating during SPS process and results in better densification [36]. Therefore, the density of Zr0.25Ti0.75B2 was increased by the addition of GNPs.  3.2. Phase analysis  Figs. 1 and 2 show the XRD patterns of the solid solution ceramics sintered at 1800 °C for 5 min. The XRD patterns of all samples contain the characteristic peaks of the (Zr,Ti)B2 solid solution (JCPDS 01-0893924) (Fig. 1). Due to their similar lattice type and atomic radii difference, which is less than 15%, ZrB2 and TiB2 fulfilled the HumeRothery requirements. Using ZrB2 and TiB2, a continuous solid solution was produced, where TiB2 completely entered the ZrB2 structure via a substitution reaction. Therefore, no separate peaks of ZrB2 or TiB2 were obtained in the XRD pattern of this solid solution. Another evidence for the formation of one single-phase solid soluto higher 2θ angles tion is the peak shift (inset in Fig. 1), which indicates a decrease in the lattice parameters with the increasing amount of TiB2 in the solid solution composition. When Ti atoms with a smaller atomic radius (0.145 nm) are replaced by Zr atoms with a larger atomic radius (0.159 nm), distortion occurs, and the lattice undergoes compression, which will cause the lattice parameters to decrease. Fig. 2 shows the XRD patterns of the GNP-containing samples. The XRD results revealed that the solid solution formation was not  Fig. 1. XRD patterns of (a) Zr0.25Ti0.75B2, (b) Zr0.50Ti0.50B2, and (d) Zr0.75Ti0.25B2.  (c) Zr0.58Ti0.42B2,  26111  \\x0c', \"M.K. Akarsu and I. Akin  Fig. 2. XRD patterns of and (d) 58Zr-42Ti-G.  (a) Zr0.25Ti0.75B2,  (b) 25Zr-75Ti-G,  (c) Zr0.58Ti0.42B2,  Fig. 3. Raman spectra of (a) initial GNPs, (b) 25Zr-75Ti-G, and (c)58Zr-42Ti-G after SPS.  Zr0.25Ti0.75B2 and Zr0.58Ti0.42B2  interrupted by the addition of 1 vol% GNPs. Furthermore, no secondary phase was detected for the GNP-containing samples (Fig. 2). Moreover, the nonexistence of shoulders in the XRD patterns confirmed the homogeneous (Zr,Ti)B2 solid solution formation. The peaks of shifted to lower 2θ angles with the addition of 1 vol% GNPs (inset in Fig. 2), indicating the formation of solid solutions. Normally, when C atoms with a smaller radius (0.071 nm) than that of B atoms (0.088 nm) enter a lattice, the lattice parameters are expected to decrease. However, when C in the graphene nanoplatelet form entered the structure of the solid solutions, the peaks of the solid solutions shifted to lower angles, that is, the lattice parameters increased in accordance with Bragg's law. The increase in the lattice parameters could be related to the tensile strain of the lattice. This phenomenon is also supported by phonon softening (a red shift of 2D (or G′) and G) in the Raman analysis (Fig. 3), which indicates tensile strain in the lattice [37].  3.3. Raman analysis  Raman analysis was conducted to characterize the graphene in the structure of the samples after SPS. The most prominent feature of  Ceramics International 46 (2020) 26109-26120  monolayer graphene is the presence of G band at 1582 cm−1 and G' band at 2700 cm−1 in its Raman spectrum [16]. The G band indicates the layered structure of graphene. The spectrum of single-layer graphene presents a sharp and symmetrical G' band. The Raman spectrum of disordered graphite shows a D mode around 1350 cm−1 due to the sp2 sites. The D mode becomes active in the presence of disorder; otherwise, it does not appear in the spectrum of perfect graphite [38]. The Raman spectra of GNP-containing samples are shown in Fig. 3. For 25Zr-75Ti-G 58Zr-42Ti-G, cm−1/ and the D/G bands at 1338 1527 cm−1 and 1329 cm−1/1589 cm−1 were obtained, respectively. For 25Zr-75Ti-G, both the D and G bands were red-shifted by ~12 and ~55 cm−1, respectively. Red shift occurred due to a decrease in phonon energies caused by the stretching of the lattice because of graphene doping [39]. Additionally, the stretching of the C-C bond, symmetry breaking, and anisotropy of the graphene lattice under strain could be the other reasons for the redshift of the spectrum [40]. Due to the hexagonal nature of the graphene lattice, it has armchair edges that increase the D band intensity by elastically scattering charge carriers [39]. Therefore, an increase in the D band intensity was observed for both samples, which indicates an increment of disorder in graphene. The ID/IG ratio of 25Zr-75Ti-G and 58Zr-42Ti-G was measured as 4.39 and 1.21, respectively. The increase in the ID/IG ratio corresponds to an increase of disordered in the structure [38]. Furthermore, a decrease in the intensity and sharpness of G' bands was observed for the GNP-containing samples (Fig. 3b and c) due to the doping of GNPs [39]. Ultrasonic dispersion caused an increment in the number of layers of graphene and produced agglomerates of GNPs, which increased the disorder [41]. The increase in the thickness of epitaxial graphene causes a shift in the G' band to lower frequencies [42].  3.4. Microstructural characterization  Zr0.75Ti0.25B2  Figs. 4 and 5 show the images of the fracture and polished surfaces of the sintered samples, respectively. Intergranular and transgranular fracture modes were observed on the fracture surfaces of monolithic sintered TiB2 (Fig. 4a) and ZrB2 (Fig. 4b), respectively. Except for Zr0.25Ti0.75B2, all samples exhibited the transgranular fracture mode. Zr0.25Ti0.75B2 exhibited an intergranular fracture mode like monolithic TiB2 due to its high TiB2 content (Fig. 4c). The yellow circles shown in Figs. 4f and 5b indicate the cleavage planes of and Zr0.50Ti0.50B2, which are characteristic of the transgranular fracture mode. No secondary phase was detected in the samples, and it was proved that a continuous solid solution was achieved for all the samples with a relative density of more than 98%. Uniformly distributed pores were observed in the SEM images of monolithic samples (Fig. 4a and b). Zr0.75Ti0.25B2 had the lowest relative density (~98.6%) among all the samples. The pores were primarily round in shape and could be detected at the triple junction points in the SEM images of the polished and etched surfaces of this sample (indicated by yellow circles in Fig. 5e). The monolithic sintered TiB2 was composed of equiaxed-shaped grains (Fig. S2) with an average size of ~19.7 μm. However, excessive grain growth was observed for monolithic sintered ZrB2. The sample contained large equiaxed grains with an average size of ~57.2 μm and straight grain boundaries (Fig. S2). The grain growth is very sensitive to temperature. The melting temperature of ZrB2 is lower than that of TiB2; therefore, the excessive grain growth of ZrB2 could be related to the higher SPS temperature (1800 °C) used herein. Fig. 5 shows the SEM images of the chemically etched surfaces of samples. The amount of pores was minimal. Finer and more homogeneous microstructures were achieved via the solid solution formation as compared to those of monolithic ZrB2 and TiB2. Except for Zr0.5Ti0.5B2, the grain size increased with the increasing ZrB2 content in the sample structure. The average grain sizes for and Zr0.75Ti0.25B2 were ~12.8, ~9.5,  Zr0.50Ti0.50B2,  Zr0.58Ti0.42B2  Zr0.25Ti0.75B2,  26112  \\x0c\", 'M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  Fig. 4. SEM images of the fracture surfaces of (a) monolithic sintered TiB2, (b) monolithic sintered ZrB2, (c) Zr0.25Ti0.75B2, (d) Zr0.50Ti0.50B2, (e) Zr0.58Ti0.42B2, and (f)  Zr0.75Ti0.25B2.  ~15.4, and ~17.8 μm, respectively (Table 1). Fig. 6 shows the SEM images of the chemically etched surfaces of the GNP-containing samples. The color differences in the microstructures were due to the topographic differences between grains. GNPs were detected at the triple junction points along the grain boundaries (Fig. 6d). The images of the microstructures of samples revealed that the presence of GNPs increased the grain size of the samples. The average grain sizes for 25Zr-75Ti-G and 58Zr-42Ti-G were ~15.2 and ~27.1 μm, respectively. Normally, it is expected that the wrapping effect of GNPs should inhibit grain growth. However, the agglomeration of GNPs at the grain boundaries, which was also detected by Raman analysis, might have reduced the average number of pinning points at grain boundaries and caused an increment in grain size [43].  3.5. Mechanical properties  The Vickers hardness and indentation fracture toughness of the samples are presented in Table 1. The average hardness of monolithic sintered ZrB2 and TiB2 was measured as ~13.1 and ~22.7 GPa,  respectively. Although the microstructure of monolithic TiB2 comprises large grains and a high relative porosity (~4.7%), the hardness of the sample is in agreement with the previously reported results [3,41,44]. However, the measured hardness of ZrB2 was much lower than the available theoretical data (~23 GPa) [3,44,45]. This is probably due to excessive grain growth of monolithic ZrB2 during SPS and consistent with the relationship between the grain size and hardness of a ceramic material. The solid solution samples exhibited higher Vickers hardness than pure ZrB2, even if the Ti content was small. The Ti(Zr) atoms are replaced by the Zr(Ti) atoms during solid solution formation. Owing to this substitutional replacement of atoms, lattice distortion occurs. Depending on the size of the atoms, this lattice distortion creates local compressive or tensile stress fields, which improve the mechanical properties [41]. The shift observed in the XRD patterns (Figs. 1 and 2) is an indication of the lattice distortion, as explained in section 3.2. Moreover, the degree of lattice distortion increases with the increasing a/c ratio of hexagonal crystal systems [46]. Consequently, herein, three possible reasons are considered for the increase in the Vickers hardness with the formation of solid solutions: lattice distortion, high theoretical hardness of TiB2, and fine grain size. The incorporation of TiB2 in a  26113  \\x0c', 'M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  Fig. 5. SEM images of the polished and chemically etched surfaces of magnification image of Zr0.75Ti0.25B2.  (a) Zr0.25Ti0.75B2,  (b) Zr0.50Ti0.50B2,  (c) Zr0.58Ti0.42B2, and (d) Zr0.75Ti0.25B2 and (e) high  small amount into ZrB2 (for Zr0.75Ti0.25B2) increased the hardness of ZrB2 by ~42%. The calculated a/c ratio for this sample is ~0.8926, which is the lowest ratio among those of all the solid solution samples. The maximum hardness of ~26.3 GPa was achieved for Zr0.25Ti0.75B2, which has the highest a/c ratio (~0.9057) for solid solutions without GNPs, a fine grain size (~12.8 μm), and a high densification (~98.9%). The addition of GNPs to 25Zr-75Ti and 58Zr-42Ti reduced the hardness by ~10% and ~12% for 25Zr-75Ti-G and 58Zr-42Ti-G, respectively. The indentation method was used to compare the fracture toughness of the samples (Table 1), which showed a different trend than the hardness. The fracture toughness of monolithic sintered ZrB2 and TiB2 was measured as ~3.08 and ~3.93 MPa⋅m1/2, respectively. Except for Zr0.25Ti0.75B2, 25Zr-75Ti-G, and 58Zr-42Ti-G, lower fracture toughness was obtained for the other samples as compared to the case of monolithic borides. Fig. 7 shows the crack paths generated by the Vickers indentation tests. Crack deflection mechanism was detected in Zr0.25Ti0.75B2. Moreover, the cracks propagated along the boundaries of small grains (1-2 μm), and branched (Fig. 7a). These two toughening mechanisms, i.e. high lattice distortion and refinement of grain size  μm), (~12.8 could be the possible reasons for the high fracture (~4.93 MPa⋅m1/2) of Zr0.25Ti0.75B2. toughness In the crack deflection mechanism, fracture toughness is enhanced by the decreasing energy release rate and stress intensity factor at the crack tip [47]. Crack deflection was also observed (Fig. 7b) for Zr0.58Ti0.42B2, which has a hardness of ~24 GPa and low fracture toughness of 2.7 MPa⋅m1/2. However, the crack path was straight without significant bending until the deflection point was reached, as indicated with a yellow arrow in Fig. 7b. Zr0.5Ti0.5B2 had the second highest hardness of ~24.4 GPa, which could be related to its fine grain size (~9.5 μm) and high densification (~99%). Some previously reported hardness values for this composition are lower than those obtained herein. Inagaki et al. [20] have reported a hardness of ~20 GPa and the relative density of ~95% for In another study, a similar hardness of ~21.8 GPa was obtained using hot pressing at 1975 °C [44]. However, herein, a fraclow as ~3 MPa⋅m1/2 was obtained for Zr0.5Ti0.5B2. ture toughness as The main reason for such a low toughness could be the presence of cleavage planes, which are characteristic for the brittle transgranular fracture mode (indicated by a yellow circle in Fig. 5b). Similarly,  Zr0.5Ti0.5B2.  26114  \\x0c', \"M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  Fig. 6. SEM images of the polished and chemically etched surfaces of (a) 25Zr-75Ti-G, (b) high-magnification image of 25Zr-75Ti-G, (c) 58Zr-42Ti-G, and (d) highmagnification image of 58Zr-42Ti-G. Note: the magnifications are different. The arrow and circles indicate GNPs.  Zr0.75Ti0.25B2  Zr0.75Ti0.25B2  had the lowest hardness (~18.7 GPa) and fracture toughness (~2.7 MPa⋅m1/2) among all the samples. Bright crystalline facets with river pattern strips were detected on the Zr0.75Ti0.25B2 surface, which correspond to the cleavage fracture (indicated by a yellow circle in Fig. 4f). Due to the brittle transgranular nature of the cleavage fracture, exhibits crack propagation in the in-plane opening mode (Mode I) [48]. The interface between ceramic and graphene is important to interpret the mechanical properties of the samples [21]. Fig. 7(c-f) show the SEM images of the fracture surfaces and crack paths created by the in 25Zr-75Ti-G and 58Zr-42Ti-G. The GNP Vickers indentation tests addition did not change the fracture mode of the ceramics. 25Zr-75Ti-G exhibited an intergranular fracture mode (Fig. 7c), while 58Zr-42Ti-G possessed a transgranular fracture mode (Fig. 7d). It was revealed that the addition of GNPs to 25Zr-75Ti activated effective toughening mechanisms such as bending (with a 90° angle), corrugation, pull-out, and crack deflection (Fig. 7c). Owing to their exceptional in-plane strength, which acts as a barrier to crack propagation, GNPs are effective crack deflectors [49]. The pull-out structure consumed energy and resisted the widening and propagation of cracks by the interaction between the embedded reinforcement and the matrix [50]. An improvement in fracture toughness was achieved because of the energy dissipating nature of corrugated structures and the energy requirement of the bending structure, which increased the damping of cracks [36]. Additionally, ~500 nm overlapping of the GNPs was detected at the grain boundaries of 58Zr-42Ti-G (Fig. 7d). Overlapping induces higher crack resistance by the efficient network of high-surface-area GNPs as a result of the superior platelet-to-platelet contact [49]. The straight path of the indentation crack in 58Zr-42Ti-G did not significantly deviate; howfor 58Zr-42Ti-G ever, the bridging of cracks by GNPs was obvious (Fig. 7f). Although GNP addition decreased the hardness of the samples by 10-12%, the fracture toughness of the samples was increased by 2-12% via the effective toughening mechanisms. Probably, GNP addition contributes to the increase in toughness in two ways; i) some GNPs  enter the solid solution systems, inducing lattice distortion, and ii) some GNPs stay in the individual form and increase the toughness by the abovementioned toughening mechanisms. However, Raman analysis (Fig. 3) revealed that the G' band was broadened and its intensity decreased. This means that the layered structure of GNPs was diminished because of ultrasonic dispersion. If the GNPs had preserved their layered structure in this study, a higher fracture toughness would have been achieved. Ocak et al. [51] have reported that upon adding GNPs to the (Zr,Ti)C structure, the hardness decreased and fracture toughness increased. Competing effects of densification, GNP-induced strengthening, grain growth, and interfacial reactions on the hardening of the ceramics can be the reason for the lower hardness of GNP-containing samples [35]. Improved mechanical properties were also reported by Mondal et al. [13] for spark plasma sintered ZrB2-TiB2 (5-30 wt%) solid solution ceramics. The authors reported the highest Vickers hardness and fracture toughness of ~20.6 GPa and ~4.7 MPa⋅m1/2 for ZrB2-30 wt% TiB2 with a densification of ~97%, respectively. Moreover, Li et al. [52] have reported that the solid solution reaction, which occurs at the interface by replacing of Ti2+ in the center of the titanium ring with Zr2+, increases the internal energy and reduces the movement speed of grain boundaries during the SPS; this causes a reduction in the grain size. Therefore, the mechanical properties of the samples improved. Karthiselva et al. [53] have also reported a significant improvement in mechanical properties by solid solution formation. The authors [53] achieved the highest nanohardness and indentation fracture toughness of 34 GPa and 3.9 MPa⋅m1/2, respectively, for Zr0.5Ti0.5B2. This superior μm. Grain hardness could be related to the fine grain size, ~1.5 boundary area increment by the reduction of grain size causes shortdistance movement of any dislocation until the dislocation reaches a grain boundary. In this way, the strength can be enhanced.  26115  \\x0c\", 'M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  Fig. 7. SEM images of the indentation crack paths on the polished surfaces of (a) Zr0.25Ti0.75B2, and (b) Zr0.58Ti0.42B2; fracture surfaces of (c) 25Zr-75Ti-G and (d) 58Zr-42Ti-G; and indentation crack paths on the polished surfaces of (e) 25Zr-75Ti-G and (f) 58Zr-42Ti-G. Note: the magnifications are different.  3.6. Oxidation behavior  To investigate the oxidation behavior of samples, TG analysis was conducted in the temperature range of 25-1100 °C (Fig. S3). Mass change was not observed until 800 °C, and then, a mass change of 1.17, 1.42, 2.25 and 2.62% was noticed for ZrB2, TiB2, 25Zr-75Ti-G, and 25Zr-75Ti, respectively. The furnace oxidation experiments, conducted at 1100 °C for 180 min, confirmed the TG analysis results. A thick, white oxide layer formed on the surface of the samples after they were exposed to oxidation at 1100 and 1200 °C for 180 min. The samples retained their structural integrity after the oxidation tests. Specific mass change per unit area for all the samples at 1100 and 1200 °C is presented in Table 2. Monolithic ZrB2 exhibited the best oxidation performance with a smallest weight change among all the samples after oxidation at 1100 and 1200 °C for 180 min. The oxidation regime of monolithic ZrB2 is very similar to that of TiB2; however, ZrB2 has a better oxidation behavior. When the temperature is above 800 °C, ZrB2 starts to oxidize according to the following equation. [54]:  attributed to the strong covalent nature of ZrB2 [55] and formation of ZrO2 and liquid B2O3 according to Eq. (4). The minor weight change measured for monolithic ZrB2 (Table 2) is an indication of the low oxidation rate of ZrB2. Additionally, the specific mass change values obtained after the oxidation of samples at 1200 °C were higher than those acquired at 1100 °C for all the samples. For atmospheric re-entry systems in which flight trajectory is exposed to severe conditions, active oxidation is a concern [56]. The inward diffusion of oxidant species caused oxidation reactions that occur at the refractory substrate-scale interface. Evaporation-resistant refractory oxide formation owing to the oxidation of the metal in diboride produces a porous skeleton. The porous skeleton fills with glassy boric oxide formed by the oxidation of boron [57]. The passive oxidation of ZrB2 and TiB2 forms a protective oxide film on the surface of the samples. The weight change of monolithic TiB2 is strongly dependent on the oxidation temperature. When TiB2 is exposed to air at high temperatures (> 800 °C), the following reactions are known to occur [41,54]:  ZrB  2 ( s  )  +  5/2 O  2 ( g  )  →  ZrO  2 ( s  )  +  B O  2  3 ( ) l  (4)  TiB  2 ( s  )  +  5/2 O  2 ( g  )  →  TiO  2 ( s  )  +  B O  2  3 ( ) l  (5)  The better oxidation performance of ZrB2 than that of TiB2 could be  The parabolic weight gain behavior of TiB2 (Fig. S3) indicates that  26116  \\x0c', 'M.K. Akarsu and I. Akin  Table 2  Ceramics International 46 (2020) 26109-26120  Main composition and thickness of  the oxide layers, and specific mass change of the samples after oxidation at 1100 °C and 1200 °C for 180 min.  Sample  Composition and the thickness of  the oxide layer  Total thickness of  the oxide layer (μm)  Specific mass change (mg/cm2)  1100 °C  1200 °C  1100 °C  1200 °C  1100 °C  1200 °C  TiB2 25Zr-75Ti  25Zr-75Ti-G  ZrB2  TiO2 + B2O3 1) TiO2 layer, ~35 μm 2) Ti-depleted ZrO2 layer, ~30 μm 3) ZrO2-TiO2 layer, ~65 μm 1) TiO2 layer, ~25 μm 2) Ti-depleted ZrO2 layer, ~15 μm 3) ZrO2-TiO2 layer, ~60 μm ZrO2 + B2O3  TiO2 + B2O3 Ti-depleted (~8 μm) ZrO2-TiO2 layer  ~60 ~130  Ti-depleted (~5 μm) ZrO2-TiO2 layer  ~100  ZrO2 + B2O3  ~40  ~95 ~220  ~200  ~65  8.34 12.43  11.83  4.38  12.82 24.71  20.93  5.95  oxidation occurs according to Eq. (5). The liquid B2O3 formed on the outer surface acts as a barrier to oxygen diffusion. Moreover, a weight loss may occur when the temperature is above ~1200 °C due to the rapid evaporation of B2O3. Then, the oxidation of TiB2 occurs without the formation of the liquid/glassy B2O3 layer on the surface. Furthermore, the oxidation behavior of the (Zr,Ti)B2 solid solutions can be considered as the summation of the oxidation reactions of TiB2 and ZrB2 according to the following equation:  ( Zr  1  x  −  Ti  x  )  B  2 ( s  )  +  5/2 O  2 ( g  )  →  (1  −  x ZrO )  2 ( s  )  +  xTiO  2 ( S  )  +  B O  2  3 ( / l  g  )  (6)  The X-ray diffraction patterns obtained for the oxidized sample surfaces revealed the formation of a combination of rutile TiO2 (JCPDS 00-004-0551), B2O3 (JCPDS 00-006-0297) and ZrO2 (03-065-2357) [58]. The formation of these compounds is controlled by Eq. (6). Most probably, crystalline B2O3 was formed during cooling after the oxidation test. Fig. 8 shows the cross-sectional morphology and energy dispersive X-ray spectroscopy (EDS) maps of the oxide scale formed on 25Zr-75Ti after oxidation at 1100 °C for 180 min. A continuous oxide scale can be observed on the sample surface, as shown in Fig. 8a. The maps of titanium, zirconium, and oxygen showed that the oxide scale had a thickness of ~130 μm and comprised of three layers: (1) a porous rutile layer with a thickness of ~35 μm, TiO2 outer (2) a Ti-depleted ZrO2 layer with a thickness of ~30 μm (indicated by the yellow dotted lines in Fig. 8b), and (3) a ZrO2-TiO2 layer with a thickness of ~65 μm. The maps of B and O revealed that B2O3 remained in the oxide scale; however, B was difficult to detect by EDS. Moreover, it was observed  the  from that of  the oxide scale was different  that the morphology of original sample. The formation of the Ti-depleted layer during oxidation has also been reported by Desmaison et al. [59] for equimolar TiB2-AlN composites. Herein, the formation of the Ti-depleted layer is probably related to the diffusion of Ti into the sample surface during oxidation at 1100 and 1200 °C. Drouelle et al. [60] have recently prepared finegrained and coarse-grained Ti3AlC2 samples using SPS and hot isostatic pressing, respectively. The authors evaluated the relationship between the microstructure and oxidation protection and observed less protection for the coarse-grained samples after their oxidation at 800-1000 °C due to the formation of a TiO2 + Al2O3 layer, where Ti and Al-depleted grain boundaries existed. In our previous study [41], the oxidation behavior of TiB2 was reported, and the EDS analysis of the oxidized surface (~60 μm thick) revealed the formation of crystalline rutile TiO2 and B2O3 phases after the oxidation of TiB2 at 1100 °C for 180 min. The formation of the Ti-depleted layer was not observed for monolithic TiB2 after its oxidation at 1100 and 1200 °C. Fig. 9 shows the cross-sectional morphology and EDS maps of the oxide scale formed on 25Zr-75Ti-G after oxidation at 1100 °C after 180 min. For the GNP-containing samples, oxidation resulted in the formation of three oxide layers, i.e. a porous TiO2 layer with an average thickness of ~25 μm, a Ti-depleted ZrO2 layer with a thickness of 15 μm, and a TiO2-ZrO2 layer with a thickness of ~60 μm. The reduction in the thickness of the Ti-depleted layer on the GNP-containing samples is an indication of reduced oxygen penetration through the oxide scale [61]. The outermost layer was not continuous, and the  Fig. 8. Cross-sectional images of (a) 25Zr-75Ti, and the elemental composition maps of (b) titanium (The yellow dotted lines indicate Ti-depleted layer), (c) zirconium, (d) oxygen, and (e) boron after oxidation at 1100 °C for 180 min. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)  26117  \\x0c', 'M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  Fig. 9. Cross-sectional images of (a) 25Zr-75Ti-G, and the elemental composition maps of (b) titanium (The yellow dotted lines indicate Ti-depleted layer), (c) zirconium, (d) oxygen, (e) boron, and (f) carbon after oxidation at 1100 °C for 180 min. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)  thickness of the layer was not homogeneous. The maps of Ti and O proved that the outer layer was mainly TiO2. C and B were detected in these oxide layers. The amount of oxygen gradually decreased towards the bottom layer, and an unoxidized region was also observed. The thickness of the total oxide layer and mass change value for 25Zr-75TiG were smaller than those for 25Zr-75Ti. Upon increasing the oxidation temperature from 1100 to 1200 °C, the oxide layer thickness of 25Zr-75Ti and 25Zr-75Ti-G increased from ~130 (Fig. 8a) to ~220 μm (Fig. S4) and from ~100 (Fig. 9a) to 200 μm (Fig. S5), respectively. The formation of a thin Ti-depleted layer was observed after oxidation at 1200 °C (Figs. S4 and S5). Additionally, the position of the Ti-depleted layer shifted to the intermediate region after oxidation at 1200 °C. The outer layer after oxidation at 1200 °C resulted in a higher amount of porosity than after oxidation at 1100 °C probably due to the evaporation of B2O3. The reduced specific mass changes for the GNP-containing samples are in agreement with the TG analysis results (Fig. S3). This might be an indication of the enhancement in the oxidation resistance of sample with the addition of GNPs [62]. The completely oxidized samples consisted of an oxide layer and unoxidized region. The mixed oxides TiO2 and ZrO2 and C were detected in the oxide layer of GNP-containing samples (Figs. 9 and S5). The rapid oxidation of ZrB2 and TiB2 might have induced very rapid formation of oxide layers and trapping of C in this region. The thickness of the outer oxide layer of Zr0.25Ti0.75B2 decreased from ~130 to ~100 μm (Figs. 8 and 9) when the GNP-containing sample was oxidized at 1100 °C for 180 min. A similar decrease in the oxide layer thickness from ~220 to ~200 μm was observed after the oxidation of the GNP-containing sample at 1200 °C for 180 min. This decrease in the thickness of the oxide layers of GNP-containing samples is an indicator of the enhancement of the oxidation resistance of the solid solutions. Furthermore, Nieto et al. [62] have stated that GNP addition enhanced the oxidation resistance of TaC-GNP composites to an extreme oxidizing environment. GNP wrapping is the reason for the enhancement in oxidation resistance due to grain sealing, which inhibits the influx of oxygen through grain boundaries (Fig. 7c) [62]. The oxidation resistance of borides is highly dependent on the protective capabilities, structures and melting temperatures of oxide scales formed on the surface of the samples exposed to air. The oxidation of TiB2 is dominant during the oxidation test of the solid solutions since there is more TiO2 in the sample and TiB2 oxidizes faster than ZrB2. This caused the formation of TiO2 in the outermost layer after the oxidation of the solid solution at 1100 and 1200 °C. Moreover, the specific mass change (Table 2) of ZrB2 and TiB2 was lower than that of Zr0.25Ti0.75B2 and 25Zr-75Ti-G after their oxidation at 1100 and 1200 °C. This could be attributed to the presence of the Ti depleted layer as explained above. The poor oxidation resistance of solid solutions disagrees with the observations previously reported by Zhang et al. [63]. The authors prepared spark plasma sintered TaC-HfC solid solutions and reported an improved oxidation resistance of these solutions than those of pure TaC and HfC. They attributed the improved oxidation behavior to the formation of the Ta2Hf6O17 phase, and the structure of HfO2 was filled with molten Ta2O5. In the abovementioned study, high-temperature plasma flow was used for the oxidation test at the temperature of ~2700 °C, which is high enough to melt Ta2O5 (whose melting point is 1850 °C). The formed oxide scale filled with the glassy phase is reported as an ideal structure for high-temperature applications [61,63]. TiO2 has a lower melting point (~1843 °C) [64] than that of ZrO2 (~2715 °C) [64]. The oxidation temperatures (1100 and 1200 °C) used herein were not high enough to melt TiO2, and the protective multiphase layer formation was not observed. Therefore, studies on the oxidation behavior of (Zr,Ti)B2 solid solutions under extreme conditions are needed. Most probably, at higher oxidation temperatures, the solid solutions would have better oxidation performance due to the formation of molten TiO2 or a possible multiphase between ZrO2 and TiO2 [64]. Furthermore, the binary B2O3 systems with group IV transition metal oxides (such as TiO2, ZrO2, and HfO2) tend to show microphase separation or immiscibility [61]. The system having phase separation is characterized by an increased liquidus temperature and viscosity, which suppress the evaporation of B2O3. Higher viscosity leads to a lower rate of oxygen diffusion through the oxide scale to the unaffected zone [65]. The tendency for phase separation is proportional to cation field strength (ratio of valence to ionic radius), and the presence of high-field cations promotes the immiscibility of the B2O3 glass surface and improves the oxidation protection. The reported [66] field strengths for Ti4+ and Zr4+ are 1093 and 772 nm−2, respectively. Herein, the oxidation behavior of Zr0.25Ti0.75B2 was tested at 1100 and 1200 °C, and phase separation (immiscibility) was not detected. A solid solution sample having higher ZrB2 content might exhibit immiscibility and better oxidation performance with the substitution of the higherfield cation Ti4+, as reported by Talmy et al. [66] for ZrB2-SiC ceramics with 10 mol% TiB2.  4. Conclusions  Herein, (ZrxTi1-x)B2 solid solution ceramics (x = 0, 0.25, 0.50, 0.58, 0.75, 1) were produced by SPS at 1800 °C under 50 MPa for 5 min TiB2 completely entered the ZrB2 structure and produced a complete solid solution for all compositions. Zr0.25Ti0.75B2 exhibited the best densification (~98.9%) and mechanical properties (~26.3 GPa and ~4.9 MPa⋅m1/2). However, monolithic TiB2 and ZrB2 showed better  26118  \\x0c', \"M.K. Akarsu and I. Akin  Ceramics International 46 (2020) 26109-26120  oxidation behavior than Zr0.25Ti0.75B2. The poor oxidation resistance of the solid solutions could be attributed to the formation of a Ti-depleted layer after their oxidation at 1100 and 1200 °C. GNP addition contributed to the improvement of toughness by creating lattice distortions and activating the toughening mechanisms such as pullout, crack bridging, and crack deflection. Moreover, GNP addition enhanced the oxidation behavior of Zr0.25Ti0.75B2. GNPs can potentially be used to improve the mechanical performance of boride solid solutions; however, the layered structure of GNPs should be preserved after their addition to the sample.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgments  The authors acknowledge Istanbul Technical University for financial support through project number: BAP-42163. Moreover, Melis Kaplan Akarsu is thankful to the Scientific and Technological Research Council of Turkey (TUBITAK) and Council of High Education (YOK) for 2211/C Domestic Doctoral Scholarship and 100/2000 Ph.D. Scholarship, respectively. Authors are greatly indebted to Prof. Dr. G. Goller for providing production and characterization facilities and thankful to H. Sezer, B. Yavas, and MSGSU Merlab (Dr. O. Ormanci) for their assistance with microstructural investigations, SPS experiments, and Raman analysis, respectively.  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2020.07.106 .  References  [7]  [6]  [4]  [1]  L. Brewer, H. Haraldsen, The thermodynamic stability of refractory borides, J. Electrochem. Soc. 102 (1955) 399-406. [2] C.L. Yeh, H.J. Wang, A comparative study on combustion synthesis of Ta-B comInt. 37 (2011) 1569-1573, https://doi.org/10.1016/j.ceramint. pounds, Ceram. 2011.01.024. [3] W.G. 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  "PDF": "Mechanical properties of ZrB2- and HfB2-based ultra-high temperature ceramics fabricated by spark plasma sintering.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 33 (2013) 1373-1386  Mechanical properties of ZrB2and HfB2-based ultra-high temperature ceramics fabricated by spark plasma sintering  E. Zapata-Solvas a,∗  , D.D. Jayaseelan a , H.T. Lin b , P. Brown c , W.E. Lee a  a Centre for Advanced Structural Ceramics, Imperial College London, SW7 2AZ, UK b Materials Science and Technology Division, Oak Ridge National Lab, TN 37831-6068, USA c Dstl, Porton Down, Salisbury, Wiltshire SP4 0JQ, UK  Received 1 November 2012; received in revised form 7 December 2012; accepted 10 December 2012  Abstract        ±  ±  Flexural strengths at  room  temperature, at 1400 C  in air and at  room  temperature after 1 h oxidation at 1400 C were determined  for ZrB2 and HfB2 -based ultra-high temperature ceramics (UHTCs). Defects caused by electrical discharge machining (EDM) lowered measured strengths significantly and were used to calculate fracture toughness via a fracture mechanics approach. ZrB2 with 20 vol.% SiC had room temperature strength of 700   90 MPa, fracture  toughness of 6.4   0.6 MPa, Vickers hardness at 9.8 N  load of 21.1   0.6 GPa, 1400 C strength of 400   30 MPa and room  temperature strength after 1 h oxidation at 1400 C of 678   15 MPa with an oxide  layer  thickness of 45  \\u242em. HfB2 with 20 vol.% SiC showed  room  temperature strength of 620   50 MPa,  fracture  toughness of 5.0   0.4 MPa, Vickers hardness at 9.8 N  load of 27.0   0.6 GPa, 1400 C strength of 590   150 MPa and room temperature strength after 1 h oxidation at 1400 C of 660   25 MPa with an oxide layer thickness of 12  \\u242em. 2 wt.% La2O3 addition to UHTCs slightly reduced mechanical performance while increasing tolerance to property degradation after oxidation and effectively aided  internal stress relaxation during spark plasma sintering (SPS) cooling, as quantified by X-ray diffraction (XRD). Slow crack growth was suggested as the failure mechanism at high temperatures as a consequence of sharp cracks formation during oxidation. © 2013 Elsevier Ltd. All rights reserved.  ±  ±     ±     ±   1   ±     ±  ±  ±  ±  ±      5   Keywords: Spark plasma sintering (SPS); Electrical discharge machining (EDM); Flexural strength; Fracture surface; High temperature; Oxidation  1.   Introduction  ZrB2 and HfB2 are used  as baseline materials  for ultrahigh  temperature ceramics (UHTCs), which are promising for aerospace  applications due  to  their  exceptional  combination of high melting point, high  thermal conductivity and excellent mechanical properties1-10 summarized  in Table 1. These features make UHTCs particularly attractive and potential candidates for use  in sharp wing  leading edges (WLEs) and nose tips for thermal protection of hypersonic vehicles during re-entry conditions. However, the surface temperature on the stagnation region can potentially exceed 2000 C under an aggressive oxidizing flow environment, which limits the range of materials that     ∗  Corresponding author. Current address: Instituto de Ciencia de Materiales de  Sevilla (CSIC-Universidad de Sevilla), C/Américo Vespucio 49, 41092 Seville,  Spain. Tel.: +34 654471922; fax: +34 954552870.  E-mail addresses: ezapata@us.es, ezapata@icmse.csic.es  (E. Zapata-Solvas).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2012.12.009     can be used to those that are able to retain their dimensional and structural  integrity. Rapid heating of UHTCs during  re-entry can produce  thermal gradients  that  induce  thermal stresses as 1500 high as 400 MPa at  C  in  leading edge components.11 The magnitude of  the  thermal stresses depends on  the  thermal expansion and Young modulus mismatch between  the different phases contained  in  the UHTC composites,  the higher  the thermal expansion mismatch, the bigger the thermal stresses. In addition, it has been demonstrated that ZrB2 -based UHTCs plastically deform at  temperatures  lower  than 2000 C,12,13 which aids  relaxing  thermal stresses.  It  is believed  that  the higher a materials room temperature strength, the higher its high temperature strength will be. However,  the stress needed  to produce plastic deformation at high  temperatures with  the same strain law14 according  rate  is proportional  to a grain size power  to the creep equation, which  is  the converse. In addition, Hu and Wang showed  that by  lowering grain size  the flow stress and the  temperature, at which  the plastic deformation  is  induced, are  reduced.13 In  fact, high  strength at  room  temperature  is           \\x0c', '1374   Table 1  E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386  Summary of some physical properties of monolithic ZrB2 and HfB2 . Coefficient of thermal expansion values are from RT to 600 C.     Property  Density (g/cm3 )   Young’s modulus (GPa)   Poisson’s ratio   −1 )  Coefficient of thermal expansion (K  RT Electrical conductivity (S/m) RT thermal conductivity (W/mK1 )   Melting temperature (  C)      ZrB2  HfB2  6.1191 4896 0.116  ×   10  −6 1  5.9  107 1 609 324510  11.2121 4807 0.128  6.3   9.1  1046 338010  −6 1  ×   10 × 106 1        ×  ×  −1  −1  to 629 MPa1,6,14,15 but   increased lowering either matrix or matrix reinforcement grain size as a consequence of flaw size reduction,1,2 which has driven researchers  to carry out  investigations with  the goal of getting the highest strength at room  temperature.1,2,7,9,15-21 A balance between microstructural parameters such as grain size and second phases is required to maintain strength at high temperature as high as possible in order to maximize structural stability during performance, even in cases where plastic deformation takes place. Monolithic ZrB2 has a flexural strength ranging from 275 this can be  improved  to a maximum value of 1089 MPa9 by adding SiC,  typically  from 10 vol.% to 30 vol.%. SiC grain size has a strong  influence on flexural strength and with micro SiC particles  (1-6  \\u242em), values  ranging from 700 MPa to 1100 MPa can be obtained for ZrB2 -SiC composites.1,7,9,15,17-21 However, owing  to SiC’s  lower coef−6 K ficient of  thermal expansion, CTE  (4.7   10 from RT −6 K to 300 C)22 compared  to ZrB2 (5.9   10 from RT  to 600 C),1 a balance is needed between thermal expansion mismatch and SiC content and high SiC concentrations (>30 vol.%) are not beneficial for ZrB2 -UHTCs. HfB2 -based ceramics follow which are  the same  trend as ZrB2 -based ceramics  in  terms of flexural strength.1 There have been fewer studies of high  temperature strength or mechanical properties  than  in ZrB2 -based ceramics, presumably due  to higher HfB2 ’s cost and density (Table 1). However,  the  combination of HfB2 ’s higher density, which enables change  to  the centre of mass of hypersonic vehicles enhancing manoeuvrability during re-entry,11 in addition to higher oxidation resistance make HfB2 -based UHTCs an attractive alternative  to ZrB2 -based UHTCs. A key  to  the performance of these materials at high temperatures is the balance between  the different mechanisms  involved during mechanical loading. Subcritical or slow crack growth has been established as  responsible  for  fracture of UHTCs  in flexion at high  temperatures   1500 C),23 but still  too  low  to produce plastic deformation. Flexural strength at high temperatures is enhanced by SiC13,24 and MoSi2 25 additions; however there have been few experimental studies of UHTC high temperature strength due in large part  to  the practical difficulties of mechanical  testing at temperatures >1600 C. UHTCs high electrical conductivity (107 S/m)9 makes them suitable for shaping by electrical discharge machining (EDM), which enables production of complex shapes otherwise impossible  to obtain by diamond cutting  tools. Nonetheless reduced flexural  strength  has  been  reported  in UHTCs  composites machined by EDM26 that can potentially limit the applicability  ≤  (T         of the components because of residual surface cracks formed on machining. Optimal machining procedures for ceramics allow rapid, cost-effective material removal while minimizing or controlling the residual damage. Flexural strength is flaw controlled through Irwin’s equation27 ;  σ  √  = KIC Y a  (1)  σ  where  is flexural strength, KIC is critical stress  intensity factor or fracture toughness in Model I, Y is stress intensity factor, which depends on flaw size and shape, and a is flaw size, which in surface cracks corresponds to flaw depth. Surface cracks are mainly  the strength  limiting flaws observed  in other materials with machining damage. Y depends on the geometry of the flaw, an expression for calculating  it was reported by Newman and Raju28 for a sample loaded in bending containing semi-elliptical flaws, which  is a  reasonable approximation  to  surface flaws observed in ceramics. It is the goal of this work to characterize the mechanical properties of ZrB2 and HfB2 -based ceramics and  the  influence of second phase additions such as SiC and La2O3 , on their mechanical properties at room temperature and at 1400 C. In addition, a preliminary study of room temperature retained strength after 1 h oxidation at 1400 C has been carried out, giving some  insights about possible  fracture mechanisms at 1400 C. Furthermore, a study of the size and shape of EDM induced flaws and their effect on mechanical properties, with the aim of developing a machining framework in which no damage is introduced into and onto UHTCs, is presented.           2. Experimental details  2.1. Sintering procedure  ρ  ρ  ρ  d50   ZrB2 powder (>99%, d50   2.5  \\u242em,   = 6.085 g/cm3 , Sigma Aldrich, Gillingham, UK), HfB2 powder (>99%, d50   5.0  \\u242em, ρ = 10.5 g/cm3 , Sigma Aldrich, Gillingham, UK), SiC powder (␣-SiC,  \\u242em,   = 3.217 g/cm3 , Good  Chemicals, Huntingdon, UK) and La2O3 (>99%, d50  99%,   0.7  Fellow \\u242em,  10   = 6.51 g/cm3 ,  Fluka  chemicals  supplied  through  Sigma Aldrich, Steinheim, Germany) were  used  to  form  different ␣-SiC (6H) have a hexagUHTC compositions. ZrB2 , HfB2 and  lattice parameters of a = 3.17 ˚A, c = 3.53 ˚A, onal structure with  a = 3.139 ˚A,  c = 3.473 ˚A  a = 3.08 ˚A,  c = 15.12 ˚A,  and  respectively. Powders with the following compositions were prepared, ZrB2 , ZrB2+ 20 vol.% SiC  (hereafter  termed ZS20), ZrB2+ 20 vol.%  SiC + 2 wt.%  La2O3 (hereafter  termed  ZS20La), HfB2 , HfB2+ 20 vol.% SiC  (hereafter  termed HS20), HfB2+ 20 vol.% SiC + 2 wt.% La2O3 (hereafter termed HS20La). They were wet ball milled  in plastic  containers using ZrO2 balls in ethanol  for nearly 24 h and dried using a  rotary evaporator under vacuum. Then,  they were  sintered  in  a  spark plasma sintering (SPS) furnace (Type HP D 25, FCT systeme GmbH, Rauenstein, Germany). A 40 mm graphite die was used and covered with graphite felt  to reduce heat  loss,  the  temperature was monitored  and  controlled  by  an  optical  pyrometer. A sintering temperature of 2000 C was used for monolithic ZrB2 , HS20 and HS20La, 1950 C for ZS20 and ZS20La and 2100 C               \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   Young’s modulus, and  the Poisson’s ratio. Then  stress can be calculated from Eq. (2) as:  ν  (cid:6)  (cid:7)   (cid:2)  (cid:3)  =  σ  E  +  v  1   ∂ε  ∂(sin2 ψ)  1375  the residual  (3)     for HfB2 , a heating rate of 100 C/min was maintained with an applied load of 50 MPa for all compositions and holding times between 3 and 6 min.  In addition, a  load  rate of 50 MPa/min was applied at 1300 C during SPS.     2.2. Materials characterisation  2.2.1. Microstructural analysis and residual stress measurement  Bulk density was determined using Archimedes’ method with water  as  the  immersing medium. Relative density was calculated by dividing measured density by  theoretical density calculated using the rule of mixtures. Geometrical methods were also used  to confirm  the validity of measured densities. The phase composition was analyzed with X-ray diffraction (XRD) in a Philips PW7100 diffractometer using Cu-K␣ radiation. All compositions were characterised by scanning electron microscopy  (SEM). Specimens  for SEM were prepared using conventional methods involving successive steps of grinding and polishing with diamond slurries and cloths embedded with up to 1  \\u242em diameter particles. The  specimens were observed  in an SEM equipped with a field emission gun  (LEO 15, JEOL, Tokyo, Japan). Samples were examined  in secondary electron (SE) imaging mode for fractography and backscattered electron (BS) image mode for atomic number contrast. Energy dispersive X-ray  spectroscopy  (EDS) was used  for phase  identification. Residual  stress was measured using an X-ray diffractometer (PW7100, Philips, Eindhoven, The Netherlands) following  the sin2 ψ  method devised by Cohen et al.29,30 High angle XRD peak shifts were obtained for 2θ  angles  larger  than 90 in  the presence of residual stress.29 The residual stress in different UHTC the sin2 ψ compositions was  thus measured using   XRD  technique at the same angle, focusing on the near-surface region.30 In  this study,  the (3 0 2)  lattice planes of ZrB2 and HfB2 were chosen as standards due to their high angle and intensity, located at 143.047 and 146.616 , respectively, according to ICDD 340423 and 38-1398, respectively. Residual stress  is expected  to be  localized at SiC boundaries with ZrB2 or HfB2 . However, ZrB2 and HfB2 peaks were chosen  to study  residual stress  in UHTC composites due  to  the absence of SiC XRD peaks at angles higher  than 90 and  the  low  intensity of SiC peaks as a consequence of  its  low volume  fraction  (20%). Stress was calculated using  the variation of  interplanar spacing within  the crystal, calculated from the XRD pattern. In this method lattice strains were measured at different  tilt angles (ψ). The surface stress component was determined using a least-squares line for the lattice strain as a function of sin2 ψ . If a biaxial stress exists in the surface layer, the lattice strain is correlated with the stress as follows30 :              = dφψ − d0  ε  d0  =  (cid:2)  (cid:3)  1   v  +  E  σφ sin2 ψ  −  (σ11 +  σ22 )   (2)  (cid:4) v  (cid:5)  E  where d0 is the interplanar spacing of the stress-free state, dφψ the lattice spacing in the direction defined by   and  σ φ is the stress component along  the direction   defined  in  the surface layer,  is  the  tilt angle,  and  are  the  stress components along the principal directions on the wear surface, E is the  σ 22  σ 11  ψ ,   ψ  φ  φ  (∂ε/∂(sin2 ψ)) was measured by determining the slopes of the plots of residual strain versus sin2 ψ  for each UHTC composition, which is the same slope as in a plot of lattice spacing dφψ versus sin2 ψ .  ×  ×  ×  ×  ×  ×  2.2.2. Mechanical properties  × 5 mm thick) billets were cut using EDM SPSed (40 mm dia  to obtain 25 mm   2.2 mm   1.7 mm bars, all bar surfaces were ground using a 1200 grit diamond platen to remove surface damage up to 25 mm   2.05 mm   1.55 mm dimensions. Finally, the tensile face was polished to 1  \\u242em and all edges were chamfered. Final dimensions of  the bars used  for flexure  strength characterization were 25 mm   2.0 mm   1.5 mm. Microhardness (HV1.0) was evaluated by Vickers  indentation, using a 9.81 N load for 15 s. Reported values were calculated from an average of 10  indentations on a single specimen of each composition, estimating projected area of  indentor prints after  indentation. Flexural strength was calculated  in air using a 3-point bending fixture with a span  length of 20 mm  in a conventional deformation machine  (Zwick/Roll, Munich, Germany) operated at cross-head speed of 0.5 mm/min according  to ASTM C 1495 standard. At least five bars were tested for each composition to calculate average flexural strength. Flexural strength in 3-point bending was calculated according to fracture mechanics single beam theory27 using the following equation;  σ  = 3FL  2w2h  (4)  σ  L   where  is the flexural strength, F is the applied force at failure, is  the span  length, w  is  the sample width and h  is  the sample height. Flexural strength  is  the maximum  tensile stress  in the sample and  is  located exactly opposite  to  the central point of contact, which  is not usually  the breakage point where surface defects produce the mechanical failure. As a result, flexural strength  is usually overestimated  in all cases and a correction using  fracture mechanics  is needed  in order  to compare  these values with those obtained using another loading configuration, such as 4-point bending  in which  the field stress  is homogeneous between the two inner points of contact and no correction is needed. According to fracture mechanics single beam theory approximation,24 which is reasonable for bend bar dimensions of current ASTM standards, the stress variation between the central contact point and the outer points is linear and a correction can be easily evaluated by measuring  the  fracture origin. Flexural strength in air at 1400 C was determined using a 3-point bending fixture with a span  length of 20 mm  in a custom-designed pneumatic loading system equipped with a furnace operated at stressing rate of 30 MPa/s according to ASTM C 1211 standard. Three bars per composition were  tested  to calculate average flexural strength.                           \\x0c', '1376   E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386  2.2.3. Oxidation treatments  An assessment of  room  temperature flexural strength after oxidation was carried out. Because  the  fracture surfaces after testing  in air at 1400 C were completely oxidized, 3 samples of each composition were oxidized for 1 h at 1400 C and then tested at room temperature with the aim of identifying the flaws produced during oxidation and understanding  the mechanical failure at high  temperature. Oxidation  time of 1 h was chosen because  it  is  the approximate  time  it  takes  to stabilize  temperature in the furnace and carry out the test at 1400 C, therefore similar oxidation and flaws are expected  in both experiments. Mass gain per unit area was measured after oxidation in all bars tested. Cross sections of oxidized samples were cut, ground and polished to 1  \\u242em and oxide layer thickness measured in the SEM at 10 random locations on the cross section and the average value used as the oxide layer thickness.           2.2.4. Fracture surfaces  Fracture  surfaces of broken bend bars were characterised using a stereoptical microscope (Leica Microsystems, Munich, Germany) for preliminary characterisation of the main features, followed by detailed SEM characterisation and examination  in a scanning coherence  interferometry (SCI) microscope (Leica Microsystems, Munich, Germany) for accurate measurement of the flaw size that produced the bending failure. Stress intensity factors were calculated  from  the flaw depth and  length using Newman and Raju28 equations as described by Quinn31 for the case of semi-elliptical flaw shape.  Ydepth = M√  Q  H2  √  π  (5)  Ysurface = M√  Q  H2  √  π  (6)  where Ydepth and Ysurface are   the stress  intensity  factors at  the deepest point and at the edge of the sample, respectively, M, Q, H2 , H1 and S are geometric terms:  M   =  (cid:8)  1.13   −   0.09  (cid:4) a  c  (cid:5)(cid:9)  +  (cid:6)  −0.54   +  0.89  0.2   +   (a/c)  (cid:7) (cid:4) a (cid:5)2 (cid:7) (cid:4) a (cid:5)4  h  +  (cid:6)  0.5   −  1  0.65   +   (a/c)  +   14  (cid:4)  1   −  (cid:4) a  c  (cid:5)(cid:5)24  c  (7)  H1 =   1   −  (cid:4)  0.34   +   0.11  (cid:4) a  c  (cid:5)(cid:5)   (cid:4) a  h  (cid:5)  (8)  H2 =   1   −  (cid:4) (cid:6)  1.22   +   0.12  (cid:4) a (cid:4) a  c  (cid:5)(cid:5)   (cid:4) a  h  (cid:5)  +  0.55   −   1.05  c  (cid:5)0.75 +   0.47  (cid:4) a  c  (cid:5)1.5  (cid:7) (cid:4) a  h  (cid:5)2  (9)  Φ  =  (cid:10)  Q   =  (cid:11)  0+1.464  (cid:4) a (cid:3) (cid:11)  c  (cid:5)1.65  for   a/c  ≤   1   (10)  S  =  (cid:2)  1.1   +   0.35  (cid:4) a  h  (cid:5)2  a  c  (11)  where a is the flaw depth, c is the half length of the flaw and h is the height of the sample as in Eq. (4). Eqs. (5) and (6) do not include an additional term fw in the original Newman-Raju analysis since, for nearly all cases involving ceramics, the flaw width is small relative  to  the component width and  fw is 1.0. The  (a/c))24 term for M is genuine. The Newman-Raju Y factors (1  in Eqs. (5) and (6) have been widely used and are  included  in several ASTM standards including C1421 for fracture toughness of ceramics, E 740 for metals and ISO standard 18756. Experience and subsequent analyses have shown  these factors  to be reliable and accurate. Y factor is affected by the flaw shape and varies around the flaw periphery. A flaw goes critical when one portion of the flaw reaches KIC , so the maximum Y is of primary concern and responsible for fracture. An assessment based on fracture mechanics was carried out to calculate fracture  toughness KIC using Eq. (1) and Newman and Raju stress intensity factors from the flexure strength measurement and characterisation of  the flaws’ depth and  length. The latter assessment allowed estimation of flexural strength of UHTCs without any machining damage,  just considering flaw sizes according to current processing routines.  2c   −  3. Results  All samples attained high density values, higher than 99% of theoretical density (TD) except ZrB2 which was 95.8%. Theoretical density was calculated using the rule of mixtures. Table 2 shows SPS conditions and density values measured after sintering. Phase analysis by XRD confirmed all  starting phases remained in UHTCs after SPS. Fig. 1 shows secondary electron (SE) images of ground, polished and etched UHTCs surfaces after SPS. ZrB2 are the larger \\u242em), finer dark grains are SiC (1-2  \\u242em) and grey grains (2-6  light grey grains are La2O3 (1  \\u242em) difficult to distinguish from ZrB2 grains, which are, respectively,  labelled with 1, 2, and 3 in Fig. 1c as they were confirmed with EDS. Homogeneous dispersion of SiC  in ZS20 and SiC/La2O3 in ZS20La  is found  in the ZrB2 matrix. SiC and La2O3 are usually  located  together forming  small agglomerates as an  interconnected network  in the ZrB2 and HfB2 matrix. However,  less  than 5% of SiC and La2O3 particles were detected separately. On the other hand, SiC agglomerates as  large as 20-30  \\u242em were found  in HfB2 -based composites as seen in Fig. 1b indicated by an arrow. ZrB2 (3 0 2) peak shift at 143 as a consequence of residual stress is illustrated in Fig. 2a. Fig. 2b shows lattice spacing dφψ versus sin2 ψ  in ZrB2 -based UHTCs. In addition, a similar study of  residual stress was carried out  in HfB2 -based UHTCs and results of residual stress calculated as described in Section 2.2.1 are shown  in Table 3. Residual stresses were calculated using Eqs.  (2) and  (3).  In addition, Young’s modulus and Poisson’s ratio values substituted  in Eqs. (2) and (3) for residual stresses calculation were  taken from  literature as  listed  in Table 1. It  is noticeable  that ZrB2 and HfB2 ceramic matrices are under  the effect of compressive internal stresses. Table 3 summarizes the flexural strength measured  in 3 point bending and calculated using Eq. (1), hardness (HV1.0), matrix grain size and residual stress  in matrix grain boundaries of UHTCs. Flexural strength               \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   1377  Table 2  SPS conditions of ZrB2 -based ceramics and bulk density, TD was calculated using the rule of mixtures with the data given in Section 2.1.  Sintering temperature (     C)   Holding time (min)  Pressure (MPa)   Bulk density (g/cm3 )   % of TD  ZrB2 ZS20  2000   6   50   5.83   0.04  5.46 ±  0.04  5.51   0.03   ±  95.8  99.1 ± 99.6   ±   0.7  1950  4  50   0.7  ZS20La  1950  3  50  ± ± ± ±  ± ± ± ±   0.5  HfB2 HS20   2100   5   50   10.42    0.5   99.2    0.5  2000   4   50   9.03    0.03   99.9    0.3  HS20La   2000   3   50   8.96    0.04   99.7    0.4  Fig. 1. SEI of: (a) ZS20, (b) HS20, (c) ZS20La, and (d) HS20La.  0,8119 -0,1  0,0  0,1  0,2  0,3  0,4  0,5  0,6  0,7  0,8120  0,8121  0,8122  0,8123  0,8124  d  s  a p  c  i  g n  (  A  )  sin  2 ψ   ZS20   ZS20La  141  14  2  14  3  14  4  14  5  5000  10000  15000  20000  25000  30000  35000  40000  45000  50000  I  n  t  n e  s  i  t  y  (  a  r  b  .  n u  i  t  s  )  2θ (  °)   Tilt=-50.77   Tilt=-42.13   Tilt=-33.21   Tilt=-22.79   Tilt=-0  Tilt=22   Tilt=33   Tilt=42  .13  Tilt=50  .77  .79  .21  a  b  Fig. 2. XRD analysis of threshold stress in ZrB2 -based UHTCs; (a) peak shift at high angle in ZS20 and (b) d spacing versus peak shift in ZrB2 -based UHTCs.  Table 3  Flexure strength in 3 point bend test corrected using a fracture mechanics approach, threshold stress, hardness (HV1.0) and ZrB2 grain size in UHTCs.  UHTCs   RT flexural strength (MPa)   Residual stress (MPa)   HV 1.0 (GPa)   Grain size (\\u242em) 10 4.0   ZrB2 ZS20   450   ± ± ± ± ±   40   -   16.5   ± ± ± ± ± ±   0.9   700    90   240   ± ±   50   21.1    0.6   ± ±   0.5  ZS20La   600    70   0    40   19.3    0.6   3.5  12 6.2    0.4  HfB2 HS20   510    50   -   19.8    0.7   620   50  690 ±  40   130   ± ±   40   27.0    0.6   ± ±   0.4  HS20La  40    40   24.2    0.8   5.5    0.5                  \\x0c', '1378   E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386  Fig. 3. Fracture surface analysis of ZS20La sample with flexural strength of 603 MPa: (a) optical micrograph, (b) high magnification optical micrograph, (c) SEM  image, (d) height map by SCI, and (e) height profile along white arrows in (d).  has been corrected as described in Section 2.2.2. SiC and La2O3 phase additions  increase flexural strength and hardness, whilst decreasing matrix grain size. The matrix grain size reduction is a consequence of the lower temperatures and shorter times used during SPS, as shown  in Table 2, owing  to  the role of second phases as sintering aids. Fine SiC (1  \\u242em) addition allows the system to fully densify and control ZrB2 or HfB2 grain growth in UHTC composites under less severe sintering conditions than monolithic UHTCs, as was pointed out by Monteverde24 for hot pressed samples and confirmed  in  this study for SPS samples. ZrB2 and HfB2 grain sizes are approximate values determined from fracture surface observations. Fracture surface characterization was carried out by stereoptical microscopy, SCI and SEM  to  locate  the fracture origin and measure flaw size and shape. A  typical example  is given in Fig. 3, where macroscopic characterization to locate all fracture surface  features, such a hackle, mist, mirror and  fracture origin, was carried out by stereoptical microscopy (Fig. 3a and b). Once the fracture origin location is known the flaw size and  shape were measured from  the depth profiles acquired by SCI (Fig. 3d and e). Alternatively, flaw size can be measured as well by SEM (Fig. 3c). However, it is more difficult to determine the exact flaw size without depth profile information. Fig. 3e shows a flaw depth a of 41  \\u242em, a flaw half  length c of 80  \\u242em in a ZS20La sample with a height of 1.34 mm, which failed at 603 MPa or 580 MPa after correction. Using these data with Newman and Raju’s equation gives a maximum stress intensity  factor Y of 1.51 corresponding  to a  fracture  toughness of 5.53 MPa/m1/2 . Table 4  shows  fracture  toughness KIC values and an estimate of  the flexural strength   without machining damage. A semi-circular flaw with a diameter of 20  \\u242em was assured with a stress  intensity factor Y of 1.3, which  is a reasonable flaw size and shape  in high strength ceramics due  to processing. Monolithic UHTCs strength  is  in agreement with the highest values obtained by others,1 the highest strength of 1 GPa was obtained  in ZS20 and ZS20La UHTCs and  lower strengths, around 100-200 MPa less, were obtained in HS20 and HS20La. In addition, the flaw depth and length ranges (Table 4)  σ  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   1379  Table 4  Flaw size range induced by EDM, fracture toughness KIC and estimated flexural strength   σ   of ZrB2 and HfB2 -based UHTCs at room temperature.  UHTCs  ZrB2 ZS20  ZS20La  HfB2 HS20   HS20La   Flaw depth range (\\u242em)   Flaw length range (\\u242em)   Fracture toughness (MPa/m1/2 )   Estimated flexural strength (MPa)  40-50   30-65  33-55  35-60   35-55   40-55   100-150   100-300  150-300  100-200   150-300   100-250   ±  ± ± ±  3.6   0.3  6.4 ± 5.2 ±  0.6   0.5  3.5   0.4   5.0   4.4    0.4    0.4   ±  ± ± ± ±  620   50 1070 ±  110 930   90  600   850   760    60   90   80           are much higher than any microstructural features observed indicating the significance of damage introduced by EDM. The same trend is observed in fracture toughness values, where monolithic UHTC fracture toughnesses are in agreement with literature values and ZrB2 -based UHTCs have higher fracture toughness than HfB2 -based UHTCs.1 Flexural strengths at 1400 C and at room temperature after oxidation at 1400 C  for 1 h are shown  in Table 5 as well as measured and estimated room temperature flexural strengths. No strength data at 1400 C is given for monolithic ZrB2 and HfB2 because  their poor oxidation  resistance dramatically  reduces their mechanical performance. Fig. 4  illustrates  the flexural strength at  room  temperature and 1400 C  to aid comparison. Mass gain per unit area and oxide  layer  thickness after oxidation  treatment  is also shown  in Table 5  to  reveal  the effect of oxidation. Oxide  layer  thicknesses for monolithic ZrB2 and HfB2 , ZS20, ZS20La, HS20 and HS20La are shown  in Fig. 5. HfB2 -based UHTCs have better oxidation resistance than ZrB2 based UHTCs in terms of mass gain and oxide layer thickness, 70% while monolithic ZrB2 has poor oxidation resistance with  of the volume completely oxidized. A  fracture  surface of ZS20 after 1 h oxidation at 1400 C (Fig. 6)  reveals  the  fracture mirror  in  the unaffected  (nonoxidized) volume with multiple flaws  50  (indicated by an arrow in Fig. 6a). A sharp crack  \\u242em long is identified as the critical flaw in Fig. 6b. In addition, sharp cracks of the same length        Fig. 4. Flexural strength at room temperature and 1400     C of all UHTCs.  perpendicular to the interface between unaffected and oxidized layers are seen in Fig. 6a highlighted by the ellipse.  4. Discussion              XRD analysis  indicated all  starting materials  remained  in UHTC composites after SPS and new phases were not formed during  sintering. Monolithic UHTCs  had  the  largest  grains (Table 3) owing  to  the higher  temperature needed  for SPS, 2000 C  for ZrB2 instead of 1950 C  required  to sinter ZS20 and ZS20La, 2100 C  for HfB2 instead of 2000 C  for HS20 and HS20La, as well as  the  longer dwelling  time. In addition, ZS20, ZS20La and HS20, HS20La had similar matrix grain sizes \\u242em and 6  \\u242em) because of the similar temp/time profile used in (4  processing. With regard to the secondary phase dispersion, hard agglomerates of SiC and La2O3 were not observed  in ZrB2 based UHTCs. However,  in La-doped UHTCs SiC and La2O3 particles tended to remain in close proximity as expected when two particles of similar but small size are mixed with a larger one. The presence of SiC hard agglomerates (20-30  \\u242em)  in HfB2 based UHTCs  is attributed  to  the density difference between HfB2 (10.5 g/cm3 ) and SiC (3.217 g/cm3 ) which is double than in ZrB2 -based UHTCs  and makes SiC float  in  the  ethanol solution during  thermo-rotary drying. Use of higher  solution viscosity or addition of a dispersant  to avoid  the presence of strong agglomerates after sintering is worthy of investigation. Residual  stress analysis  in  the ZrB2 and HfB2 in Zr and Hf-based UHTCs revealed the existence of an internal stress as a consequence of fast cooling during SPS. Furthermore, SiC is mechanically stressed under compression due to lower CTE val−6 compared to  6  ues than the matrix phase (4.7   10  10 of ZrB2 and HfB2 ).  In addition,  internal  stresses as high as 240 and 120 MPa were found  in ZS20 and HS20, respectively. Moreover,  the addition of La2O3 was observed  to effectively reduce  internal  stress, which  could  potentially  control  load transfer between different phases in UHTC composites altering crack  deflection  and  subsequent  fracture  toughness. Raman scattering and neutron diffraction experiments have  revealed 800 MPa in ZrB2-30 vol.% the presence of internal stresses of  3  2  SiC,32 with particle size of  for ZrB2 and  for SiC. Differences in the residual stress values measured with the XRD technique of this study are attributed to differences in SiC addition (20 compared  to 30 vol.%) as well as different matrix grain size  (4 compared  to 3  \\u242em).  It  is  intuitive  that a higher SiC concentration  increases residual stresses. In addition, SiC agglomerates  reduce  the number of matrix grain boundaries mechanically  stressed  in  contact with  SiC  as well  as  the  \\u242em   \\u242em   −6  ×  ×  \\x0c', '1380   E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386  Table 5  Room temperature measured and estimated flexural strength, flexural strength at 1400     C, flexural strength, mass gain per unit area and oxide layer thickness after  oxidation at 1400     C for 1 h of ZrB2 -based UHTCs.  UHTCs   RT flexural  strength (MPa) 450 ±  40  700 ±  90  600   70   RT estimated flexural  strength (MPa) 620 ±  50  1070   110   Flexural strength  1400     C (MPa)  RT flexural strength after  oxidation 1400 190 ±  10  620   45      C 1 h (MPa)  Mass gain after oxidation C 1 h(mg/cm2 )  1400     Oxide layer thickness C 1 h (\\u242em)  1400     ZrB2 ZS20  -  11.70   ± ± ± ± ± ±   0.48   291   ± ± ± ± ± ±   5  ± ± ± ± ±  400   ± ±   30   ± ± ± ± ±  4.54    0.21   45    5  ZS20La   ± ± ± ±  930    90   358    6   720    10   2.18    0.11   28    3  HfB2 HS20   510    50   600    60   -   290    15   8.80    0.35   248    6  620    50   850    70   590   ± ±   150   660    25   0.60    0.03   12    1  HS20La   690    40   760    50   480    40   610    30   1.97    0.10   36    3  Fig. 5. SEM pictures of UHTC cross sections after 1 h oxidation at 1400     C: (a) ZrB2 , (b) HfB2 , (c) ZS20, (d) HS20, (e) ZS20La and (f) HS20La.  Fig. 6. Fracture surface of ZS20 after 1 h oxidation at 1400     C: (a) low magnification optical image and (b) optical image of the mirror with the flaw.  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   1381  (i)   (ii)  (iii)     larger matrix grain size allowing more matrix volume  to relax contact-induced stresses, leading to reduced residual stress. This argument is able to explain differences in internal stress between ZS20 and HS20. Moreover, herein  the  internal stress has been measured in ZrB2 instead of in the SiC particles,32 which might reduce  internal  stress because not all  interfaces  in ZrB2 are mechanically  stressed as  they would be  in SiC.  In addition, XRD analyses free surface particles which are not constrained as  in  the volume and some  load could have been  released.  In addition XRD peak shift is insensitive to stresses normal to the sample surface. Some studies33,34 have detected  intragranular crack formation after rapid cooling in the SPS arising from CTE mismatch between different phases  in UHTCs. Nonetheless, the cooling ramp was not controlled allowing  the SPS  to cool down by  its own  thermal  inertia which ensures cooling  rates higher than 100 C/min. Therefore, there is no time to relax all stresses  in  the UHTCs on cooling so  inducing microcracks  in the final composites. In  this study, no cracks were detected by SEM  inspection  in all  the cooling ramp controlled UHTCs. In addition, nearly 50% of samples prepared without controlling the  cooling  ramp  showed  cracks,  lower  flexural  strength, hardness  and  fracture  toughness  than monolithic UHTCs.  It should be possible  to modify  internal stresses and subsequent load transfer between different phases in UHTCs by controlling the cooling rate with the aim of tailoring their mechanical properties. However, beyond  the  limit  that produces  intragranular crack formation the reinforcement effect is no longer beneficial. Flexural  strength  is  clearly  increased by  addition of  sec700 MPa  ond phases such as SiC and La2O3 , from 450  to  in 700 MPa in HfB2 -based ZrB2 -based UHTCs and from 500 to  UHTCs. This effect  is  likely  related  to second phase particle pinning as well as grain growth reduction during sintering compared to monolithic UHTCs. It is well known that the flaw that produces mechanical  failure  in Eq.  (1)  is proportional  to  the matrix grain size1,2 ; the bigger the grain size, the bigger the flaw and the lower the flexural strength. Therefore, reducing damage introduced by EDM (surface elliptical cracks) higher strength could have been obtained, as shown by  the estimated flexural strength  in Table 4.  In addition, La2O3 doping did not negatively affect  the UHTCs flexural strength. A similar  trend has been found by adding Al2O3 , Y2O3 35 and Yb2O3 36 to UHTCs. Tables 6 and 7 provide for comparison purposes data for UHTCs in  this study and some of  the  relevant  literature data,  including fabrication procedure, matrix grain size, hardness, fracture toughness and flexural strength at room and high  temperature is  improved by addition of SiC, from 16.5  to 21.1 GPa  in ZrB2 -based UHTCs and from 19.8  to 27.0 GPa  in HfB2 -based UHTCs. La2O3 addition reduces hardness by about 10% in Zror Hf-based UHTCs, 19.3 and 24.2 GPa, respectively, compared with SiC doped UHTCs. Internal stresses are not present in ZS20La and HS20La, therefore  load  transfer and energy adsorption under contact are  less efficient during Vickers  indentation compared with ZS20 and HS20 and  reduced contact damage  tolerance  is expected and confirmed  by  experiments.  ZS20  has  the  highest  flexural strength, 700 MPa, which is still in the medium range of literature values, however it is in the high strength range if compared  if available.7,13,23-26,33-35,37-42 Hardness   HfB2 -based UHTCs  UHTCs but with  lower   reveal  the  same  trend as ZrB2 -based fracture  toughness, 5.0  for HS20 and  with UHTCs with the same ZrB2 grain size (4  \\u242em).1 Fahrenholtz 600 MPa for 4  et al’s review1 indicates flexural strengths of  \\u242em grain size ZrB2 . However, a flexural strength of 1 GPa or higher is expected when ZrB2 grain size is reduced to 1  \\u242em as a consequence of flaw size reduction. In addition, the flexural strength of HfB2 -based UHTCs  is  in  the high strength range compared with  literature values (Table 7). Moreover, machining damage introduced by EDM has been detected and characterized  in all materials using microscopy techniques (Table 4) suggesting the origin of  the medium  range values obtained  for  the measured room temperature flexural strength of UHTCs. A  fracture mechanics analysis was used  to calculate  fracture  toughness values which are high compared with values previously found for ZrB2 -based UHTCs which range from 3 to 5 MPa/m1/2 ,1,2 although  some authors used SiC whiskers, graphite flakes or Mo particles  7 MPa/m1/2 ,17,19,43 a similar value  to  improve  fracture  toughness up  to  to  the ZS20 of  this study (6.4 MPa/m1/2 ). Three factors may be responsible for the improved fracture toughness in the UHTCs studied:  is  ×  −1 )22  low ZrO2 particle contamination during milling which  lower than 1 wt.% as no ZrO2 XRD peak was observed; good dispersion of SiC as indicated by SEM observations. It  is well established  that SiC produces deflection at  the crack  tip during crack propagation  in ZrB2-SiC composites, so a more homogeneous SiC dispersion means  that more crack deflections per unit  length are produced, subsequently improving fracture toughness; the  thermal expansion mismatch between particles  in  the different  composites  can produce  internal  stresses during  rapid SPS cooling preventing a controlled  relaxation of  the  elastic  strains.  Internal  stresses  in ZrB2 -based materials  −6 K are  connected with  the  lower CTE  −6 K −1 )1 which of  SiC (4.7   10 than ZrB2 (5.9   10 promotes  tensile stresses along SiC grain boundaries and subsequently provides a more effective  load  transfer and increase of crack tip deflection or crack energy adsorption during crack propagation. In  the case of ZS20La, a  lower fracture toughness (5.2 MPa/m1/2 ) was found compared to ZS20 (6.4 MPa/m1/2 ). The origin of the different behaviour is  related  to  the  thermal expansion coefficient of La2O3 −6 K −1 ),44 which is the highest of all the com(8-10   10 posite particles, and the location of La2O3 particles in close proximity to SiC particles (Fig. 1c and d). Therefore, a more rapid La2O3 shrinkage during SPS cooling provides  free volume to SiC particles promoting a more effective relaxation of elastic strains and subsequently reducing ZS20La fracture toughness. This effect is shown in Fig. 7, in which different UHTCs structure are illustrated indicating differences in mechanical stresses after SPS controlled cooling. This trend is supported by the XRD internal stress characterization, which indicates that fracture toughness could be tailored  through  the control of  internal stresses generated by the SPS cooling ramp.  ×  ×  \\x0c', '1 3 8 2   E  .   Z  a p a  t  a   S o  l  v  a  s   e  t   a  l  .   /   J  o u  r  n a  l   o  f   t  h  e   E  u  r  o p  e  a n   C  e  r  a  m  i  c   S o  c  i  e  t  y   3 3   (  2 0 1 3  )   1 3 7 3 - 1 3 8 6  Table 6 Fabrication method, relative density, grain size, Vickers’ hardness at 1 kg load unless specified, fracture toughnessa and flexural strength at room temperature, 1200     C, 1400     C, 1500     C and 1800     C of ZrB2 -based  ceramics.  Fabrication  method  Relative density  (%)  Grain size (\\u242em)   Hardness HV1.0  (GPa)  KIC (MPa/m1/2 )  σ RT (MPa)   σ 1200    C (MPa)  σ 1400    C (MPa)  σ 1500    C (MPa)   σ 1800    C (MPa)  ZrB2 ZS20   SPS   98.0   10  4.0   16.5   ± ± ±   0.9   3.6   ± ± ±   0.3   450   ± ± ±   40   -   -   -   -  SPS   99.1   21.1    0.6   6.4    0.6   700    90   -   400   ± ±   30   -   -  ZS20La   SPS   99.6   3.5   16.0 ± 19.3  14.9 ± 15.0 ± 17.0    0.6   5.2   0.5  2.3 ±  0.2a -  600  531 ± 704 ±  46  552 ±  98   336  498   382    70   -  655 ± -  358    6   -  500 ± 333 ±  58   31  388   23   -  ZrB2 -20 MoSi2 ZrB2 -15 MoSi2 ZrB2 -20 MoSi2 ZrB2 -15 MoSi2 -2.3C25 ZrB2 -15 MoSi2 ZrB2 -10 SiC 24 ZrB2 ZrB2 ZrB2 -5 Si3N4 ZrB2 -20 SiC-4 Si3N4 ZrB2 -18.5 SiC-3.7 Si3N4 -1 Al2O3 ZrB2 -15 SiC 26 ZrB2 -15 SiC-10 HfB2 ZrB2 ZrB2 -2.3 MoSi2 ZrB2 -15 SiC-2 MoSi2 ZrB2 -15 SiC 13 ZrB2 -30 SiC 13 ZrB2 -15 SiC 13  25  PLS  99.1   2.5    0.4    17   -  -  25  HP  97.7   1.8    0.5   -  -  25  HP  98.1   1.8    0.4   -  -   -   ± ±  -  HP  99.2   1.5   ±   0.8   -   ±  -   -   479    13   -  25  SPS  98.1   1.4  3  9  10  3.0   16.2   ±   0.5   2.6   ± ±   0.3a  643   ± ± ± ± ± ±   97   632   ±   5   -   357   ± ±   48   -  HP   100   -   4.8    0.3   835    35   203  -   -   300    35   -  37  HP   100   -   -   300   -   200   ±   503  -   -  35  HP   87   8.7   ± ± ±   0.4   2.35   ± ±   0.15   350    30   310   ± ± ±   10   -   -   -  35  HP   98   13.4    0.6   3.75    0.10   600    20   240    30   -   -   -  35  HP   98   2.4   14.6    0.3   -   730    100   250    10   -   -   -  35  HP   98   2.5   14.2   ±   0.6   4.55   ±   0.10   630   ±   20   280   ±   20   -   -   -  HP  >97  3  5   17.7 ± 18.2    0.4   4.07 ± 4.1    0.03   887 ± 763    125   -  -  255   ± ±   25   -  26  HP   97.7   ±   0.5   ±   0.8   ±   73   -   -   240    20   -  7  HP   88.8  6.04 g/cm3 5.61 g/cm3  10   8.7   0.4  18.1 ±  0.4  17.7   0.4   ±  2.4   0.2  3.4 ±  0.3  4.07   0.03   ±  350   ± ± ±   30   -   210   ±   20   -   -  7  HP  5  750    160   -   -   240   ± ±   25   -  7  HP   3   ±  ±  795    105   255    25   -  HP   100   4.5   -   -   865   ± ± ±   125   -   -   -   112   ± ± ±   12b  10b  HP   100   3   -   -   705    120   -   -   -   48   HP   100   8   -   -   500    40   -   -   -   217    16  PLS, Pressureless sintering; HP, Hot pressing. a Fracture toughness calculated via VIF (see text for explanation). b Flow stress.  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   1383  Before  cooling  AŌer cooling (100 ºC/ min)  ZS20  ZS20La  ZrB2  matrix  SiC  La2O3  Therm al Expa nsion coefficients (K-1)  ZrB2 HfB2 SiC La2O3  5.9x10-6 6.2x10-6 4.7x10-6 8-10x10-6  Fig. 7.   Illustration of shrinking effect during cooling in SPS owing to thermal expansion mismatch.  4.4  for HS20La. The presence of hard SiC agglomerates and the  larger grain size of matrix particles  in HfB2 -based (6 compared to 4  \\u242em) ceramics causes a reduction of crack deflection per unit  length  and  subsequent  reduction  in  fracture  toughness. Moreover, SiC  hard  agglomerates  do  not  produce  an effective  crack deflection  tip, making  the  toughening  effect of SiC additions  to UHTCs  less efficient. UHTCs  fabricated by SPS have high hardness and  fracture  toughness compared with other techniques. However, high fracture toughness values have been  claimed  after Vickers  indentation  fracture  toughness evaluation  (VIF).38,40,42 VIF  is a non-traditional method that  uses  a Vickers  indenter  to make  a  contact  impression on  a  polished  surface,  then  the  fracture  toughness  is  evaluated  in  terms  of  crack  length,  indentation  load,  hardness, elastic modulus, indentation diagonal size and an empirical fitting constant. VIF  is also not  included  in any of  the standard methods  for  the  ceramic  fracture  toughness determination45 and  is only  recommended  for evaluating contact damage  tol2-4  erance. For example, differences of around  times have been  found  in  the evaluation of  fracture  toughness via VIF, using  the  same ceramic materials  in different  laboratories.45 Tests  included  in  standards  are presented here,  single  edge precracked beam  (SEPB), chevron notched beam  (CNB) and surface crack  in flexure  (SCF).45 Nonetheless, a method  that has grown in popularity is the single edge notched beam method, wherein a notch  is made  in  the  tensile face with a razor blade or diamond wheel. Nishida et al.46 showed  that  in Al2O3 , an anomalously high value of  fracture  toughness was obtained when  the  notch  radius was  higher  than  10  \\u242em.  In  particuto 7.6 MPa/m1/2 lar,  the fracture  toughness  increased from 3.8  when  the notch-radius  increased  from 10  to 300  \\u242em. Some UHTC studies claim high  fracture  toughness values omitting this condition17-19,39,47,48 or with a notch-radius <250  which  is not  sufficient  for  reliable measurement of  fracture toughness  in ceramics. Therefore,  standardisation of  fracture toughness measurement  in  the UHTCs community  is required for  reliable measurements  and  subsequent data  comparison.  \\u242em43,49  \\u242em   More caution should be exercised before claiming high fracture toughness in ceramics.50 Some authors33,34 have highlighted  intragranular crack formation  in UHTCs during rapid SPS cooling owing  to  thermal expansion mismatch between particles, which may be an  issue in UHTCs without machining damage. However, in our samples no damage has been observed  in  the volume of  the materials, only in the proximity of EDM machined surfaces. Damage has been detected and quantified by different microscopy techniques  (Stereoptical Microscopy, SCI  and SEM). The  latter characterisation  showed 75  \\u242em deep and 300  long dam100-150  age before  testing. After EDM machining  \\u242em of material was removed. It appears  that at  least 200  \\u242em must be removed from each face to get UHTCs free of damage and with maximum strength, assuming  that existing cracks do not grow during further removal which may well be the case in high fracture toughness ceramics. Strength is limited by the presence of defects  from processing. Therefore, a  fracture mechanics calculation considering a semi-circular flaw of 20  \\u242em, which  is a reasonable flaw size  for high strength ceramic materials, was carried out. Subsequently, high strength in UHTCs free of EDM damage has been estimated in ZrB2 and HfB2 -based UHTCs at room temperature (Table 4). Few  high  temperature  strength  data  have  been  reported, especially  in HfB2 -based UHTCs. Tables 6 and 7  reveal  that monolithic UHTCs have  the  lowest strength at  room  temperature  and high  temperature. The  addition of SiC  is  clearly beneficial for UHTCs as the room and high temperature strength increase by a factor around two. Monteverde et al.35 studied the effect of Si3N4 in ZrB2 and ZrB2-SiC ceramics which showed better  strength at  room  temperature  than ZrB2 , however  the high  temperature  strength was  lower  than 87% dense ZrB2 . This  response  is  attributed  to  the  softening of grain boundary Zr-B-N-Si-O glass  formed during hot press  sintering. Moreover, Si3N4 has a  low  thermal conductivity which makes these materials unsuitable for hypersonic applications. Hu and Wang13 reported the highest temperature strength in ZrB2 based            \\x0c', '1384   E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386  T  a  b  l  e   7  F  a  b  r  i  a c  i t  n  o  m  e  t  d o h  ,   r  e  l  a  i t  v  e   d  e  n  s  t i  y  ,   g  r  a  i  n   s  i  e z  ,   V  i  c  k  e  r  s  ’   h  a  r  n d  e  s  s  a  t   1   g  k  l  o  a  d   n u  l  e  s  s  i  s   s  p  c e  i  fi  e  d  ,   r f  c a  t  u  r  e   t  n h g u o  e  s  s  a  d  n  fl  e  u x  r  a  l   s  t  r  e  g n  t  h   a  t   r  o o  m   t  e  m  p  e  r  a  t  u  r  e  ,   0 0 4 1     C   a  d  n  0 0 5 1     C   o  f   H  f  B  2   b  a  s  e  d   e c  r  a  m  i  c  s  .  F  a  b  r  i  a c  i t  n o  m  e  t  d o h  R  e  l  a  i t  v  e   d  e  n  s  t i  y  (  %  )  G  r  a  i  n   s  i  e  z  (  \\u242e  m  )   H  a  r  n d  e  s  s  V H  1  .  0  (  G  P  a  )  K  I  C  (  M  P  a  /  m  1  /  2  )   \\u2434  R  T  (  M  P  a  )  \\u2434  0 0 4 1     C  (  M  P  a  )  \\u2434  0 0 5 1     C  (  M  P  a  )  H  f  B  2  S  P S  8 9  .  0     2  1  9 1  .  8   ±  ±  ±  0  .  7   3  .  3   ±  ±  ±   0  .  4   0  5 4  ±  ±  ±  ±  ±  ±  ±  ±  ±  ±  ±  ±  ±   0  4  -   -  H  S  0 2  S P S  9 9  .  1   2  .  9   7 2  .  0   0  .  6   5  .  0   0  .  4   0  0 7  0  9  0  9 5  ±  ±  ±   0  5 1  -  H  S  0 2  L  a  S P S  9 9  .  6   2  .  4   4 2  .  2  0  .  8   4  .  4  0  .  4   0 0 6  0  7  0  8 4  0  4  -  H  f  B  2  7 3  H  P  0 0 1  6  -   -   0  3 3  0 1  3  0  6 2  0 5  3  H  f  B  2  1  2 -  S  i  C  7  -  H  f  C  1 -  H  f  O  2  3 2  H  P  9 9  .  2        3   9 1  .  0   ±  ±   1  .  0   -   0  7 7  5  3  -   0  1 3  ±  ±  ±   5 1  H  f  B  2  0  3 -  S  i  C  3 3  S  P S  >  9  9  2   6 2  .  0   1  .  0   3  .  9   ±  ±  ±  ±  ±  ±  ±  ±  ±  ±   0  .  3   0  9 5  0  5  -   0  0 6  5 1  H  f  B  2  0  3 -  S  i  C  2  -  T  a  S  i  2  4 3  H  P   8  .  2  6  g  /  c  m  3  1  .  5   -   3  .  6   0  .  5   5  6 6  5  7  -   0  8 4  5 3  H  f  B  2  0  3 -  S  i  C  2  -  T  a  S  i  2  4 3  S  P S  8  .  6  6  g  /  c  m  3  2   -   4  .  5  6  0  .  5  0  5  6 4  5  2 2  -   -  H  f  B  2  2  -  w  t  .  %   B  4 9 3  C  8 3  S  P S  8 9  .  3   5 1  .  6   9 1  .  5   ±  ±  ±  ±  ±  ±  ±   1  .  2   4  .  1   0  .  4  a  9  6 4  3 1  3  -   -  H  f  B  2  0  2 -  S  i  C  H  P   9 9  .  1          3   9 1  .  7   0  .  4   4  .  4   0  .  6   2  5 4  5 2  3  -   -  H  f  B  2  0  2 -  S  i  C  0  1 -  A  l  N  9 3  H  P   9 9  .  6   2   8 1  .  9   0  .  4   6  .  2   0  .  3   1  9 5  0 3  3  -   -  H  f  B  2  0  2 -  S  i  C  0 4  H  P   >  5  9  2   0 2  .  2   1  .  0   (  9  4  N  )   6  .  3   0  .  6  a  3  9 9  0 9  3  -   -  H  f  B  2  0  2 -  S  i  C  1 4  H  P   >  9  9  0  1  7 1  .  0   0  .  7   (  9  4  N  )   4  .  1   0  .  2   3  5 4  6  4  -   -  H  f  B  2  2 4  H  P   0 8  .  0   2   1 1  .  5   1  .  4   (  0  .  8  9  N  )   5  .  7   0  .  3  a  -   -   -  H  f  B  2  5  1 -  w  t  .  %   T  i  S  i  2  2 4  H  P     0  0 1  4   7 2  .  4   1  .  9   (  0  .  8  9  N  )   6  .  6   0  .  2  a  -   -   -  a  F  r  c a  t  u  r  e   t  n h g u o  e  s  s  a c  l  c  u  l  a  t  e  d   v  i  a   V  I  F   (  s  e  e  t  e  x  t   f  o  r   e  p x  l  a  n  a  i t  n o  )  .  ceramics studied to date (217 MPa at 1800 C) for matrix grain size of 8  \\u242em. However, plastic deformation occurred at 1800 C for matrix grain sizes  lower  than 5  \\u242em. Elastic behaviour with relatively high strength was found with  increasing ZrB2 grain size  introducing a microstructural  tailoring factor  that controls the high  temperature mechanical response. A systematic study of  grain  size  effect  and  SiC  content  on  high  temperature strength of ZrB2 based ceramics  is required  to understand  the mechanical responses and failure mechanisms of these systems. In addition,  this  is an effect  that has been characterised  for room  temperature strength. The effect of sintering method on mechanical properties has been analysed. Pressureless sintered ZrB2-20 MoSi2 had higher strength at 1500 C than hot-pressed ZrB2-20 MoSi2 while the room temperature strengths remained similar.  In addition, hot-pressed ZrB2-15 MoSi2 and SPSed ZrB2-15 MoSi2 have  similar  room  and  high  temperature strength. This  effect  is  explained  in  terms of  the  secondary phases formed during hot-pressing and SPS of these materials, such as SiO2 and ZrO2 which soften  the grain boundaries at high  temperatures, especially  the  residual silica pockets. The addition of carbon  to  the  system avoided  the  residual  silica pocket  formation during hot-pressing  increasing substantially high  temperature  strength of ZrB2 -15 MoSi2 .  In  the present study,  the general  trend  is  that HfB2 -based UHTCs had higher flexural strength at high  temperature  than ZrB2 -based UHTCs (Tables 6 and 7). The strength at 1400 C of ZS20 and ZS20La and HS20 and HS20La is approximately twice that of ZrB2 and HfB2 ,37 respectively. The reduction of  the strength at 1400 C in  the La2O3 -doped UHTCs  could  be  attributed  (2100 to La2O3 inducing grain boundary  softening. La2O3 C) has a lower melting point than SiC (2800 C), ZrB2 (3200 C) and (3400 HfB2 C) which  induces plastic deformation at  lower temperatures. This has been observed in other ceramic systems such as Al2O3 , which has a similar melting point (2100 C) −3 s −1 in micron-sized to La2O3 and plastic deformation at 10 Al2O3 could  be  obtained  at  1400 C with  stresses  around 100 MPa.51,52 Grain boundary chemistry is affected by thermal treatment which could be beneficial for the strength of UHTCs and Fahrenholtz53 at high  temperature. Hilmas  reported  an enhancement of up to 40% of the strength of HfB2 -SiC ceramics at 1600 C using a  thermal  treatment  in Ar of 10 h at 1550 C. This effect was attributed to a grain boundary chemistry change during  the  thermal  treatment.  Further  investigation  of  this effect and strength data at higher  temperatures  is  required  to understand high  temperature mechanical  response of UHTCs. However,  this field  is challenging since some phenomena such as creep, have only been explored very recently,12 and its effect on UHTCs properties at working temperature is unknown. To help understand the failure mechanism, strength after 1 h oxidation at 1400 C was studied in all UHTCs. After oxidation, machining damage healing by silica formation during oxidation could  take place. However,  it  seems  that oxide  layer  thickness controls the healing process being effective for thicknesses \\u242em,  <30  such as  in HS20 and ZS20La or  reducing  strength at higher  thicknesses,  such  as  in ZS20, HS20La, ZrB2 and HfB2 (Table 5). This effect  is especially  intense  in monolithic UHTCs since  they possess a poorer oxidation  resistance  than                                                  \\x0c', 'E. Zapata-Solvas et al. / Journal of the European Ceramic Society 33 (2013) 1373-1386   1385     reinforced UHTCs. Comparing flexural strength after 1 h oxidation at 1400 C with estimated flexural strength, which might be real strength without machining damage, a reduction of 69%, 52%, 42%, 23%, 32% and 19%  in monolithic ZrB2 and HfB2 , ZS20, ZS20La, HS20 and HS20La,  respectively,  is observed. UHTCs with La2O3 have better tolerance to oxidation in terms of flexural  strength  after oxidation.  In  addition,  the flexural strength reduction  indicates  that a new population of flaws has been formed during oxidation and might subsequently change the failure mechanism at high temperature. Oxide layer formation  involves expansive  reaction with a volume expansion of 4-5% due to ZrO2 or HfO2 formation,10 so stress at the interface could be generated which might produce cracks. Under load at high temperatures, the cracks may grow until the critical crack size according  to Eq.  (1)  is  reached causing mechanical  failure. Slow crack growth  is suggested as a failure mechanism  in UHTCs and it has been identified as a failure mechanism in brittle ceramics, by  the UHTCs community.23 However, complete oxidation of fracture surfaces impedes direct observation of slow crack growth. The study of fracture surfaces of samples 1 h preoxidized at 1400 C, allowed the flaw population formed during oxidation to be observed. Sharp cracks instead of semi-elliptical flaws were found as shown  in Fig. 6, which could grow under mechanical  loading at high  temperatures, producing mechanical  failure via slow crack growth. An experimental study of UHTCs strength at high temperature in argon atmosphere might enable  the observation of  fracture surfaces,  revealing whether slow crack growth contributes to failure at high temperature of UHTCs. Oxidation kinetics and UHTCs strength may be correlated with oxide layer thickness and stresses involved during oxide layer formation, a further study of strength in UHTCs preoxidized at different  temperatures and  times  is underway with the aim of establishing  the relation between oxidation kinetics and strength degradation. In  previous work,  Jayaseelan  et  al.54 suggested  that  in ZS20La an oxidation resistant lanthanum zirconate coating with melting point higher  than 2000 C  formed during oxidisation which could  lead  to oxidation  resistant ZrB2-SiC UHTCs, as highlighted by Eakins et al.55 In  the current study strengths at 1400 C for ZS20La are not substantially different from those of ZS20, and also a  lower degradation of strength after oxidation has been  revealed  in La2O3 doped UHTCs.  In addition, high hardness, fracture toughness, flexural strength at room temperature and 1400 C (relative to monolithic UHTCs) were measured in SPSed UHTCs, which  represents a promising combination of high mechanical performance and oxidation  resistance  in UHTCs to be implemented in real aerospace applications dealing with temperatures around 2000 C.                 5. Conclusions  Highly dense UHTCs were  fabricated by SPS  and  their mechanical properties evaluated. Examination of  fracture surfaces revealed the machining damage induced by EDM, which was quantified  through SCI and SEM.  In addition, a  fracture mechanics assessment allowed calculation of fracture toughness and an estimation of flexure strength without EDM machining     damage of UHTCs. High hardness and fracture  toughness are obtained  in SPSed UHTCs. Differences  in  behaviour were explained  in  terms of  residual stresses generated during  rapid cooling  in SPS and second phase distribution, which could be used  to  tailor mechanical properties of UHTCs. HfB2 -based UHTCs have higher flexural  strength at 1400 C  than ZrB2 based  ceramics. However, La2O3 addition  reduces flexural strength at high  temperatures  in UHTCs as a consequence of grain boundary  softening  at  those  temperatures. A  study of strength  in pre-oxidized ZrB2 -based UHTCs  revealed  lower degradation after oxidation  in La2O3 doped UHTCs and sharp cracks formation during oxidation suggests slow crack growth is  the  high  temperature  failure mechanism.  In  conclusion, the  combination  of  high  strength, mechanical  tolerance  to oxidation, hardness and fracture toughness make La2O3 doped UHTCs promising for aerospace applications.  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},{
  "_id": 123,
  "PDF": "Mechanical, Thermal and Oxidation Behaviour of Zirconium Diboride Based Ultra-High Temperature Ceramic Composites.pdf",
  "Text": "['Key Engineering Materials Vol. 395 (2009) pp 55-68 Online available since 2008/Oct/21 at www.scientific.net © (2009) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.395.55  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (ID: 130.207.50.37, Georgia Tech Library, Atlanta, USA-10/12/14,08:00:35)   Mechanical, Thermal, and Oxidation Behaviour of Zirconium Diboride Based Ultra-high Temperature Ceramic Composites  R. Mitra1,a, S. Upender1,b, M. Mallik1,c, S. Chakraborty1,d,  and K.K. Ray1,e 1Departmemt of Metallurgical and Materials Engineering, Indian Institute of Technology,  Kharagpur -721302, West Bengal, INDIA arahul@metal.iitkgp.ernet.in, bupender.sunkari@gmail.com, cmanabmallik@yahoo.co.in, dsubrata_1980@rediffmail.com, ekkrmt@metal.iitkgp.ernet.in   Keywords: Zirconium Diboride, Ultra high temperature ceramic composites, Mechanical Property, Thermal Shock, Ablation Resistance, Oxidation Resistance.   Abstract. A comparative study has been carried out on the mechanical properties at room temperature, thermal shock and ablation resistance as well as oxidation behaviour of ZrB2-20SiC, ZrB2-20SiC-5Si3N4 and ZrB2-20ZrC-20SiC-5Si3N4 (amounts represent volume percent) composites. Fracture toughness has been determined using either three-point bend tests on single edge notch bend specimens, or by indentation technique. Addition of Si3N4 as sintering aid leads to enhancement in flexural strength and fracture toughness in the composite without ZrC. The specimens were subjected to thermal shock by quenching from temperatures in the range of 800o-1200oC to ice cold water, and to ablation by exposure to oxy-acetylene flame at 2200oC. The composite having ZrC as constituent, exhibits the highest resistance to damage due to thermal shock and ablation, while the ZrB2-SiC composite shows the least change in mass during ablation. On the other hand, thermogravimetric experiments from room temperature to 1300oC have shown that the presence of ZrC is detrimental for oxidation resistance. Hence, the constituents of the composites need to be selected on the basis of the nature of application. The results of this study show that the investigated ZrB2 based composites bear the potential for multiple use thermal protection of re-entry type space vehicles.  Introduction  The hypersonic vehicles usually have sharp leading edges and nose cones to allow maneuverability at high velocities. The leading edges and nose cones are subjected to large aero-thermal heating, which may raise the temperatures beyond 2000°C during re-entry, and consequently the materials used for these components require severe thermal protection. Refractory zirconium diboride composites belong to the class of ultra high temperature ceramics (UHTCs) due to their high melting points exceeding 3200°C and strength retention capabilities at elevated temperatures. The UHTCs based on zirconium diboride based composites have attracted strong research interest also due to their high temperature ablation and oxidation resistance, high refractoriness, high thermal conductivity and good thermal shock resistance [1,2]. Hence, the zirconium diboride based composites are candidate materials for the thermal protection tiles of sharp leading edges and nose-cones of the hypersonic vehicles.  ZrB2 is of particular interest because of its moderate density (6.09 g.cm-3) and ability to maintain yield strength of 125 MPa at 1800oC. The density of ZrB2 is lower than that of niobium-based refractory alloys normally used in aerospace structural applications, requiring thermal protection against ablation and oxidation. However, past research results [2] show that the properties of ZrB2 based composites are strongly dependent on processing. Lack of adequate \\x0c', '56  Progress in High Temperature Ceramics   oxidation resistance at high temperature has been cited as the major deterrence for the widespread use of ZrB2 [3]. The mechanical, thermal, and oxidation properties of the zirconium diborides, carbides and their composites with SiC reinforcements have been evaluated by Opeka et al. [4] and Monteverde [5]. The addition of SiC as reinforcement to ZrB2 improves the oxidation resistance, thermal shock resistance, flexural strength, and fracture toughness [6-9]. The SiC additions also aid in processing of the materials by reducing the maximum temperature required to densify the materials, and by reducing the growth of the ZrB2 grains [10].  Further reduction in the sintering temperature is achieved by addition of Si3N4 as reinforcement, which aids in densification through removal of B2O3 from the surface of ZrB2 particles [11]. Other related studies have shown that the hot pressed ZrB2-ZrC-SiC based composites exhibit excellent ablation resistance [12,13]. The present study is focused on the microstructure, mechanical properties, thermal shock and ablation resistance, as well as isothermal and non-isothermal oxidation behaviour of the hot pressed ZrB2 based composites. The role of Si3N4 and ZrC additions on the mechanical and oxidation behaviour of the ZrB2-SiC composite has been examined.   Experiments, Results and Discussions  Experimental Procedure. The experiments carried out in this study are classified as: (i) Raw Materials and Processing; and (ii) Characterization. Raw Materials and Processing. The powders of ZrB2, ZrC, SiC, and Si3N4, were obtained from H.C. Starck, Germany. The average sizes of the ZrB2, ZrC, SiC, and Si3N4 powder particles are 5.4 ± 2.2 µm, 4.8 ± 1.7 µm, 4.0 ± 1.4 µm, and 3.3 ± 1.1 µm, respectively. The purity of the powders was 99.5% or better. The powder samples of ZrB2-20 vol% SiC (henceforth referred as ZS), ZrB2-20 vol% SiC-5 vol% Si3N4 (ZSS), and ZrB2-20 vol% ZrC - 20 vol% SiC-5 vol% Si3N4 (ZZSS) were mixed at a milling speed of 250 rpm for 2 h in the Planetary Mono Mill (model Pulverisette 6, Fritsch, Germany). The powders were wrapped in grafoil, placed in a cylindrical graphite die, and subsequently hot pressed in vacuum at 2000oC for 30 min using a uniaxial pressure of 30 MPa, into pancakes of 42 mm diameter.  Characterization. The hot pressed pancakes were sectioned using an Isomet slow speed diamond saw cutter (Buehler Ltd, USA). The sectioned samples were metallographically polished, and the microstructures were examined using secondary (SE) and backscattered (BSE) electron imaging modes on a JEOL JSM 5800 scanning electron microscope (SEM) to examine the distribution of the constituent phases, the chemical compositions of which were determined using Energy Dispersive X-Ray (EDX) analysis. X-ray diffraction (XRD) using Cu Kα radiation was used to identify the phases present. The densities of the hot pressed products were measured using the Archimedes principle with water as the immersing medium. The elastic modulus of the composites was measured using the dynamic elastic property analyzer involving vibration of the sample with natural frequency [14]. The measurement of hardness was carried out using a Vickers diamond indenter operated with loads of 20 and 30 kg. Fracture toughness of the hot pressed composites was determined by three point bend testing of single edge notch bend (SENB) specimens, following the ASTM E399 standard procedure [15]. The SENB specimens having dimensions of 4 mm (thickness) X 8 mm (width) X 36 mm (length), and ~ 4 mm deep notch along the width were tested with a crosshead speed of 0.1 mm per min on a three point bend fixture with a span of 32 mm. The notches were cut using 0.3 mm thick diamond saw. The fracture toughness, K1C was determined using the equation (1):  KIC = [3PS/BW1.5].f (a/W).                                                              (1)  where P is the load, S is the span, B is the thickness, W is the width, a is the initial crack length, and f(a/W) = 3(a/W)0.5/[2{1+(2a/W)}{1-(a/W)}1.5{1.99-(a/W)(1-a/W)(2.15-3.93a/W+2.7a2/W2)}]. Further, three-point bend tests were carried out at a crosshead speed of 0.1 mm per minute, using a fixture with 40 mm span, on specimens with dimensions of 3 mm (thickness) X 4 mm (width) X 42 \\x0c', 'Key Engineering Materials Vol. 395  57   mm (length) to determine the flexural strength following the procedure described in ASTM C 1161-94 [16]. The flexural strength, σ has been calculated using the equation (2):  σ = 3PS/2WB2.                                               (2)  Selected samples of the investigated composites were subjected to thermal shock by heating at three different temperatures including 800o, 1000o, and 1200oC in air inside a vertical chamber furnace having SiC heating elements, and subsequently quenched in water. The damage to thermal shock was assessed quantitatively through the examination of changes in hardness, extension of indentation crack lengths, and indentation fracture toughness (IFT) of the composite samples after subjecting to thermal cycling. The IFT was measured using the empirical relationship provided by Niihara et al. [17]:  KIC = 0.0309(E/H)0.4(P/c1.5)        (3)  where E is the Young’s modulus, H is the hardness, P is the load, and c is the characteristic crack length. The lengths of the cracks at corners of indentations were measured before and after the thermal cycle to determine the changes in indentation fracture toughness. The procedure followed is somewhat similar to that used for the evaluation of thermal shock resistance of the ceramics and ceramic matrix composites in previous studies [18-20]. Ablation resistance was studied on the specimens having the dimensions of 6 mm (width) X 3 mm (thickness) X 30 mm (length). A neutral oxyacetylene flame with maximum temperature of 2200oC was employed to evaluate the ablation behavior of the samples, after exposures for 5, 10, and 15 min, respectively. An optical pyrometer and a thermocouple were used measure and record the time-temperature data of the front and the unattacked back surface of the specimens. The experimental set-up for the ablation experiment was designed on the basis of the procedure discussed in Refs. 12 and 13. Furthermore, the extent of damage was quantified in terms of the changes in mass, Young’s modulus and indentation fracture toughness. X-Ray diffraction (XRD), followed by energy dispersive X-ray analyses on SEM was used to identify the oxide phases present after exposure.  Finally, non-isothermal oxidation behaviour of the composites has been studied by heating at the rate of 5oC/min in a Thermo-gravimetric (TG) Analyzer (Perkin Elmer, USA). The change in mass with increase in temperature has been measured using the dry, inert Al2O3 powder as reference. Results and Discussion. The densities of the ZS, ZSS, and ZZSS composites have been found to be 5.50, 5.33, and 5.47 g.cm-3, respectively, which are 99-100% of the theoretical values calculated from the rule of mixtures. The results of the experiments related to microstructural examination, mechanical and thermal property evaluation, as well as characterization of the resistance to oxidation, thermal shock and ablation, have been discussed separately in the sections below. Based on the results, an effort has been made to correlate the material properties with the structure and composition of the different UHTCs. Microstructure. The microstructures of the ZS, ZSS, and ZZSS composites are shown in Figs. 1(a-c), respectively, and typical XRD pattern from the ZSS composite is shown in Fig. 2.  The SEM images [Figs. 1(a-c)] clearly show a uniform distribution of phases in the microstructure. The ZrB2 and ZrC phases appear bright, while the SiC and Si3N4 look dark. The identity of the major elements in each of these constituents has been confirmed by EDX analysis. Examination of the composition of the dark regions located at the grain boundary triple points, using EDX has shown evidence of Si and N enrichment. In a previous study on the ZrB2-Si3N4 composites, it has been shown that the B2O3 present as impurity in the ZrB2 is reduced by reaction with Si3N4 at around 1550oC during processing, and a liquid phase is formed, particularly at the grain boundary triple points. Such a reaction leads to the formation of BN and borosilicate type glass [11]. In the present \\x0c', '58  Progress in High Temperature Ceramics   study, the results of XRD analysis (Fig. 2) have shown the presence of Si3N4, which suggests that it exists and is possibly consumed only partially by chemical reaction during hot pressing.                              Fig. 1. SEM (SE) images showing the microstructures of: (a) ZS; (b) ZSS; and (c) ZZSS composites.  (c) (b) (a) SiC ZrB2 \\x0c', 'Key Engineering Materials Vol. 395  59   102030405060708090100010002000300040005000600070008000mQQBBBBBBBBQmSi3N4SiCBZrB2Intensity (A.U.)2θθθθ (Degree)                   Fig. 2. XRD pattern from the ZSS composite showing the peaks of ZrB2, SiC and Si3N4.  Mechanical Properties.  The elastic modulus, hardness, flexural strength and fracture toughness of the investigated composites and those of monolithic ZrB2 [21] for the purpose of comparison, are shown in the Table 1. The elastic moduli of the composites, calculated using the rule of mixtures, are also shown in Table 1 for comparison with the experimentally determined values. The experimentally determined Young’s modulus has been compared with the values predicted by the rule of mixtures in some of the previous studies [22-24] on composites, for assessing the strength and integrity of the particle-matrix interfacial bond.  A comparison of the experimental values of elastic modulus (EEXP) with those predicted by the rule of mixtures (EROM) shows the EEXP / EROM ratio to be close to 1, which suggests excellent bond integrity between the matrix and the reinforcement in case of the ZS and ZSS composites. However, the ratio of the experimentally determined Young’s modulus to that from the rule of mixtures is the least in case of the ZZSS composites, suggesting that interfacial bonding between the ZrC and the other phases present is not as strong as that in the ZS and ZSS composites.  The elastic modulus of the ZS composite with higher volume fraction of ZrB2, is higher than those of the ZSS and ZZSS. The Young’s modulus phase of ZrB2 is higher than those of ZrC and Si3N4.  On examination of the data presented in Table 1, it is clear that the ZSS composite possesses the highest values of hardness (21.2 GPa), flexural strength (520 MPa) and fracture toughness (7.8 ± 0.6 MPa√m). Again, the improvement in fracture toughness with respect to that of monolithic ZrB2 is the most significant in case of the ZSS composite (Table 1). The fracture toughness of the ZS, ZSS and ZZSS composites are higher than that of the monolithic ZrB2 by 56%, 111%, and 27%, respectively. Figs. 2(a-c) show the typical fractographs, representing the respective fracture surfaces of the ZS, ZSS, and ZZSS composites, generated through flexural tests. The fractograph of the ZS composite, depicted in Fig. 2(a), suggests a mixed intergranular and transgranular mode of failure. The features suggesting transgranular failure include grain facets and river patterns. Transgranular failure is observed in the coarser grains, while the locations containing finer grains show evidence of intergranular failure [Fig. 3(a)]. In a previous study [25] based on fracture behavior of Al2O3, it has been shown that anisotropic ceramics with fine grain size tend to exhibit primarily intergranular failure, while the transgranular mode becomes predominant with increase in the grain size. A close examination of the fracture surface [Fig. 3(a)] also indicates that the ZS composite has a wide distribution of grain sizes, which is primarily responsible for the variation observed in the mode of failure. In contrast, the fracture surfaces [Figs. 3(b) and (c)] of the ZSS and the ZZSS composites are primarily intergranular. The average size of the grains appears to be between 2 and 4 µm. The grain size distribution appears to be the most homogeneous in the ZSS composite. The observation \\x0c', '60  Progress in High Temperature Ceramics   of intergranular fracture surface for the finer and more uniform grain size is consistent with the observations recorded for the ZS composite. The higher hardness, flexural strength and fracture toughness of the ZSS composite is attributed to its finer and more or less uniform matrix grain size. It may be noted that the volume fraction of SiC and Si3N4 are exactly same in the ZSS and ZZSS composites. The lower hardness and flexural strength of the ZZSS is composite is in spite of the high hardness and flexural strength reported for ZrC [26,27]. The lower than expected hardness and flexural strength of the ZZSS composite can perhaps be explained on the basis of inadequate interphase interfacial bond-integrity, as suggested by somewhat lower value of the Young’s modulus of the composite than that predicted by the ROM (Table 1). However the values of hardness and fracture toughness of the ZZSS composite are in agreement with those reported for the ZrB2-ZrC composites prepared by spark plasma sintering [28].  Table 1. Room temperature mechanical properties of the composites. *The abbreviation “ROM” stands for “Rule of Mixtures”. The data for ZrB2 is from Ref. 21.  Material  Young’s modulus (GPa) *ROM Young’s modulus (GPa) Hardness  (GPa) Flexural Strength (MPa) Fracture Toughness (MPa√m) ZrB2 490 ---21.0 480 3.7 ZS 484 489 19.4 356 5.8 ± 0.2  ZSS 467 478 21.2 520 7.8 ± 0.6  ZZSS  427 447 17.5 407 4.7 ± 0.2   Thermal Shock Resistance.  The hardness, change in the indentation crack length, and IFT measured before and after each thermal cycle experiment, are depicted through bar charts in Figs. 3(a-c), respectively. The hardness of all the composites decreases with increasing temperature drop (∆T) during the thermal cycle [Fig. 4(a)]. But the rate of decrease in hardness of the ZZSS composite is lower than that of ZS and ZSS. The change in indentation crack length, %∆c [%∆c = (cf-ci) x 100/ci, where ci and cf are the respective crack lengths of samples before and after the thermal cycle] with ∆T is shown in Fig. 4(b). It can be seen that the %∆c increases continuously with increasing ∆T for all the composites. After thermal cycle, several micro-cracks are generated inside the composites. The cracks emanating from the tip of hardness indentation, link with the micro-cracks during its propagation, which in turn contribute to the extension in crack length. This may be considered as the reason for increase in %∆c with increase in ∆T. It is also intuitive that the micro-cracks are responsible for the degradation of hardness in the composites, subjected to temperature drop. The values of IFT are also found to degrade with increasing ∆T, as shown in Fig. 4(c). In the damaged specimens with micro-cracks, the flaw size is large and hence the critical stress intensity factor responsible for crack propagation is decreased significantly. This stress intensity factor gets superimposed with the stress intensity factor of the advancing long crack front from the indentation tip, which results in a higher crack length and in turn lowers the IFT.  \\x0c', 'Key Engineering Materials Vol. 395  61                                                       (a) (b)(c) Fig. 3. Fracture surfaces of: (a) ZS; (b) ZSS; and (c) ZZSS composite. The ZrB2 and ZrC appear bright, while the SiC and Si3N4 appear dark.  \\x0c', '62  Progress in High Temperature Ceramics   06121824(a)  1200oC1000oC800oC ∆∆∆∆T (oC)RTHardness (GPa) ZrB2-SiC ZrB2-SiC-Si3N4 ZrB2-ZrC-SiC-Si3N404812162024(b)∆∆∆∆T (oC)1200oC1000oC800oC%∆∆∆∆c   ZrB2-SiC ZrB2-SiC-Si3N4 ZrB2-ZrC-SiC-Si3N4036(c)   IFT (MPa.m0.5)1200oC1000oC800oC ∆∆∆∆T (oC)RT ZrB2-SiC ZrB2-SiC-Si3N4 ZrB2-ZrC-SiC-Si3N4                                                  Fig. 4. Bar charts showing: (a) hardness, (b) indentation crack lengths and (c) values of the IFT, before and after thermal cycling between 800o, 1000o, or 1200oC and room temperature.  \\x0c', 'Key Engineering Materials Vol. 395  63   A comparison of the degradation behavior of the composites indicates that thermal shock resistance of the ZZSS composite (with ZrC) is superior to those of ZS and ZSS composites. This may be because of the marginal decrease in co-efficient of thermal expansion (CTE) of the matrix on addition of ZrC, and reduction in the mismatch between the constituent phases. The co-efficient of thermal expansion of ZrC is 5.5 x 10-6 K-1, which is less than that of ZrB2 (6.7 x 10-6 K-1). Reduction in the CTE is expected to improve the thermal shock resistance of the composites.  Ablation Resistance. The front and back-face temperatures on the plates of the ZS, ZSS, and ZZSS composites have been recorded for 5, 10, and 15 min during heating with an oxy-acetylene flame. The typical pattern of the temperatures recorded on the front and back faces of the ZS composite during heating for 15 minutes, is depicted as time-temperature plots in Fig. 5. From examination of the results in Fig. 5, it is obvious that the front face is subjected to a temperature in the range of 2000o-2200oC, while the back face shows temperatures between 1200o-1250oC, during exposure to the oxy-acetylene flame. Thus, a temperature drop of about 800oC across its thickness of 3 mm, is recorded. During the tests, the front face in each sample has taken about 200 s, while the back face has taken 100 s to reach the stable temperatures. It may be proposed that much of the heat generated due to the oxyacetylene flame is lost to the surroundings by radiation, and only a part is being conducted to the back surface. The morphology of the front surfaces of the samples exposed to the oxy-acetylene flame for 15 min is depicted in the photograph presented in Fig. 6. Comparison of the surface characteristics clearly indicates that the oxide scale formed on the ZS composite has maintained its integrity, while there is evidence of spallation in case of the ZSS and ZZSS composites. The damage during ablation has been characterized in terms of the change in mass and Young’s modulus. The change in mass per unit surface area due to high temperature exposure is shown using bar charts in Fig. 7. The bar charts indicate that all the composites have undergone a net gain in mass due to oxidation. It is clear that for a given time period of exposure, the ZS composite has shown the least gain in mass. On the other hand, the highest gain in mass after 5 min of exposure is observed in the ZZSS composite, while that after 10 and 15 min of exposure is observed in case of the ZSS composite. The elastic moduli and indentation fracture toughness values, before and after ablation for 10 min are shown using bar charts in Fig. 8(a) and (b), respectively. The drops in the Young’s moduli, recorded after ablation of the ZS, ZSS, and ZZSS composites for 10 min, are 7.2, 8.4, and 3.3%, respectively [Fig. 8(a)]. Furthermore, ablation for 10 min has led to a noticeable decrease in the IFT of all the samples, that is, 20.8, 18.5, and 8.8% in case of the ZS, ZSS, and the ZZSS composites, respectively [Fig. 8(b)]. Comparison of the drops in Young’s modulus and indentation fracture toughness clearly indicates that the decrease in modulus and indentation fracture toughness is the least in case of the ZZSS composite. Hence it can be inferred that among the investigated composites, internal damage due to ablation is minimum in the ZZSS composite, in spite of a high mass gain exhibited by it in the first 5 min of exposure to oxy-acetylene flame. The observation of the highest ablation resistance is in agreement with that in previous studies [12,13], and is attributed to the formation of a passivating oxide scale.  Figure 9 shows the typical XRD pattern from the ZSS composite surface after ablation for 10 min, showing the peaks of ZrB2, ZrO2, SiO2 (cristobalite) and ZrSiO4, formed by oxidation. Similar peaks are observed in the XRD patterns from the ablated surfaces of the ZS and ZZSS composites. The cross-section of the damaged sub-surface location and oxide scale formed on the ZSS composite due to exposure to the oxy-acetylene flame for 10 min is shown in Fig. 10. The affected zone appears to be 40-50 µm thick. Considering that the ZSS composite is the worst affected in terms of mass gain in case of exposure to the oxy-acetylene flame for 10 and 15 min, it may be proposed that the damage is superficial.      \\x0c', '64  Progress in High Temperature Ceramics   020040060080010006008001000120014001600180020002200ZSS Back face temperature Front face temperature    Time (s)Temperature (oC)                  Fig. 5. Variation of the front and back face temperatures in the ZSS composite with time elapsed during the 15 min of exposure to the tip of the oxy-acetylene flame.                             10 mm Fig. 6. Typical photograph of ZS, ZSS, and ZZSS composite specimens ablated for 15 min. \\x0c', 'Key Engineering Materials Vol. 395  65   01234567810 minutes10 minutes15 minutes5 minutes10 minutes15 minutes15 minutes5 minutesZSZZSSZSS    Mass gain (mg/cm2)0100200300400500600(a)ZZSSZSSZSElastic Modulus (GPa) Before Ablation After Ablation   0246810(b)ZZSSZSSZSIFT (MPam1/2) Before Ablation After Ablation   Oxidation Behaviour. Non-isothermal oxidation behavior studied using thermogravimetry (TG) of the ZS, ZSS, and ZZSS composites is depicted in Fig. 11. From the plots shown in Fig. 11, it is clear that the process of oxidation characterized by mass gain, is initiated at approximately 800oC, followed by a much sharper rise at around 1200oC in case of ZS, and ZSS composites. On the other hand, the ZZSS composite has shown gain in mass at a much lower temperature of 600oC, with change in the mass gain characteristics at 800oC. The net mass gains of the ZS and ZSS composites are almost identical, while that of the ZZSS composite is much higher. The ZZSS composite differs from the other two in having ZrC as a major constituent. Hence it can be inferred that the presence of ZrC in the ZZSS composite is primarily responsible for its poor oxidation resistance.                 Fig. 7. Bar charts showing the mass gain per unit area of the ZS, ZSS, and the ZZSS composite during exposure to the oxy-acetylene flame for 10 min.                  Fig. 8. Bar charts depicting: (a) elastic modulus, and (b) IFT of the ZS, ZSS, and the ZZSS composite before and after ablation for 10 min.       \\x0c', '66  Progress in High Temperature Ceramics   2030405060708090020040060080010001200140016001800Intensity (A.U.)2θθθθ (Degree) (cid:1)(cid:1)(cid:1)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)(cid:2)ZS(cid:2)ZrO2(cid:1)ZrSiO4SiO2ZrB2                   Fig. 9. Typical XRD pattern from the ZS composite surface after ablation for 10 min, showing the peaks of ZrB2, ZrO2, SiO2 (cristobalite) and ZrSiO4.                   Summary  ZrB2-SiC (ZS), ZrB2-SiC-Si3N4 (ZSS) and ZrB2-ZrC-SiC-Si3N4 (ZZSS) composites with almost 100% of their theoretical densities, have been processed by hot pressing. The microstructures have 030060090012001500020040060080010001200  Change in Mass (µµµµg)Temperature (oC) ZrB2-SiC ZrB2-SiC-Si3N4 ZrB2-ZrC-SiC-Si3N4Fig. 10. SEM (SE) image showing the cross-section of the oxide scale formed on the ZSS composite during exposure to the tip of oxy-acetylene flame. The damaged region is shown with an arrow.  Fig. 11. Results of TG analyses: Plots showing the change in mass in the ZS, ZSS, and the ZZSS composites with temperature during heating.  \\x0c', 'Key Engineering Materials Vol. 395  67   shown a uniform distribution of the constituent phases with little porosity. The results of this study may be may be summarized as follows:  (1) Among the investigated composites, the ZSS has exhibited the most optimum combination of Young’s modulus, hardness, fracture toughness and flexural strength. (2)  The fracture surfaces bear evidence of partial intergranular and transgranular failure, with the former being more predominant in case of the ZSS and ZZSS composites with finer and relatively uniform matrix (ZrB2) grain size.  (3) Thermal shock resistance of the composites have been examined with temperature differentials of 800o, 1000o and 1200oC, followed by measurements of Young’s modulus, hardness, and extension of the indentation crack lengths, have shown the least damage in case of the ZZSS composites.  (4) Ablation studies have been carried out by exposure of the ZS, ZSS, and the ZZSS composites for periods of 5, 10, and 15 min to the oxy-acetylene flame. During the tests, the front face in each sample has taken about 200 s, while the back face has taken 100 s to reach stable temperatures in the range of 2100o-2200oC and 1200o-1300oC, respectively.  In all the composites, a temperature drop of about 800oC across the thickness of 3 mm was observed, irrespective of the time of exposure.  (5) Examination of the ablated front surfaces, and analysis of the damage using measurements of change in mass, Young’s modulus, and indentation fracture toughness, have shown that the ZZSS composite is the least prone to internal damage, while the ZS composite is the most resistant to oxidation.  (6) Non-isothermal oxidation studies involving thermogravimetry, have shown that noticeable gain in mass is initiated in the ZZSS composite at 600oC, and in the ZS and ZSS composites at 800oC.  Acknowledgement  The financial support received from the Defence Research and Development Laboratory under CARS Contract No. DRDL/23/96P/05/0173/31454 is gratefully acknowledged.  References  [1] A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, D. T. Ellerby:  Key Eng.    Mater. Vol.  264 (2004), p. 493. [2] Federick Monteverde, Stefano Guicciardi and Alida Bellosi: Mater. Sci. Eng. Vol. A346 (2003), p. 310. [3] Stanley R Levine, Elizabeth J. Opila, Michael C.Halbig, James D. 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Soc., Vol. 69(4) (1986), p. 317. [28] Tsuchida Takeshi, Yamamoto Satoshi: J. Mater. Sci., Vol. 42(3) (2007), p. 772.   \\x0c', 'Progress in High Temperature Ceramics   10.4028/www.scientific.net/KEM.395   Mechanical, Thermal and Oxidation Behaviour of Zirconium Diboride Based Ultra-High Temperature Ceramic Composites   10.4028/www.scientific.net/KEM.395.55       \\x0c']"
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  "_id": 124,
  "PDF": "Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and Zirconium Compounds.pdf",
  "Text": "['Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and Zirconium Compounds  Mark M. Opeka, a* Inna G. Talmy, a Eric J. Wuchina, a James A. Zaykoski aand Samuel J. Causey b  aNaval Surface Warfare Center, Carderock Division, West Bethesda, MD, USA  bSouthern Research Institute, Birmingham, AL, USA  Abstract  The  thermal  conductivity,  thermal  expansion,  Youngs Modulus,  flexural  strength,  and  brittle-  plastic  deformation  transition  temperature  were  determined for HfB2, HfC0(cid:1)98, HfC0(cid:1)67, and HfN0(cid:1)92 ceramics. The oxidation resistance of ceramics in the  ZrB2-ZrC-SiC system was characterized as a function of composition and processing technique. The  thermal  conductivity of HfB2 other materials by a factor of 5 at room temperature and by a factor of 2.5 at 820(cid:14)C. The transition tem exceeded that of  the  perature of HfC exhibited a strong stoichiometry dependence, decreasing from 2200(cid:14)C for HfC0(cid:1)98 to 1100(cid:14)C for HfC0(cid:1)67 ceramics. The transition temperature of HfB2 was 1100(cid:14)C. The ZrB2/ZrC/SiC ceramics were prepared from mixtures of Zr (or  ZrC), SiB4, The ceramics with ZrB2 as a predominant phase had high oxidation resistance up to 1500(cid:14)C compared to  and C using  displacement  reactions.  pure ZrB2 ZrB2/SiC molar ratio of 2 (25 vol% SiC), containing little or no ZrC, were the most oxidation resis and ZrC ceramics. The  ceramics with  tant. Published by Elsevier Science Limited.  Keywords: borides, carbides, corrosion, mechanical  properties, nitrides,  thermal conductivity,  thermal  expansion.  1 Introduction  Hafnium and zirconium-based ceramics (carbides,  borides, and nitrides) display a number of unique  properties,  including extremely high melting tem perature and hardness, as well as high thermal and  electrical conductivity and chemical stability. This  combination  of  properties make  these materials  potential candidates for a variety of high-tempera ture  structural  applications,  including  engines,  hypersonic vehicles, plasma arc electrodes, cutting  tools,  furnace  elements,  and  high  temperature  shielding. The present  investigation describes  the  thermal and mechanical behavior of HfB2, HfC, and HfN ceramics synthesized by reactive hot  pressing. Since HfC has a wide homogeneity range,  variations in the properties as a function of carbon  content were also investigated.  The ZrB2-based ceramics were as relatively oxidation-resistant non-oxide  recognized early  refrac tory  compounds. Research  on  these materials  resulted in the development of ZrB2/SiC ceramics which exhibited highly improved oxidation resis tance  over  pure  ZrB2.1 (NSWCCD, unpublished data) have  Recent  investigations  shown that  ZrB2/ZrC/SiC compositions o\\x80er further improvement of oxidation resistance at very high temperatures (2200(cid:14)C).  The present article describes the oxidation beha vior of ZrB2/ZrC/SiC ceramics at temperatures up to 1600(cid:14)C as a function of component ratio. The  ceramics were  synthesized using in-situ high-tem perature  solid-state  chemical  (displacement)  reac tions.  The  advantage  of  the  method  is  the  possibility to create materials with novel and con trolled microstructures, with high chemical  com patibility of  in-situ-formed individual phases, and  phase distribution uniformity.  2 Experimental Procedure  The hafnium-based compounds were hot-pressed (75 mm diameter(cid:2)15 mm thick billets) using \\\\37 7325  mesh powders from Cerac, Inc. (Table 1). Hafnium  hydride (HfH2) was added to all starting mixtures to decrease processing temperatures. The non-stoi chiometric  HfC0(cid:1)67 according to the reaction:  ceramics  were  prepared  Journal of the European Ceramic Society 19 (1999) 2405-2414  Published by Elsevier Science Limited  Printed in Great Britain  P I I :  S 0 9 5 5 2 2 1 9 ( 9 9 ) 0 0 1 2 9 6  0955-2219/99/$ see front matter  2405  *To whom correspondence  should be  addressed. Fax: +1 301-227-4732; e-mail: opekamm@nswccd.navy.mil  \\x0c', '67HfC (cid:135) 33HfH2 ! 100HfC0(cid:1)67 (cid:135) 33H2  The ZrB2/ZrC/SiC ceramics were synthesized from 99.7% purity, \\\\37 the starting mixtures 7325 mesh) or ZrC (Cerac, 99.5% consisting of Zr (Cerac, purity, \\\\37 7325 mesh), SiB4 (Cerac, 98.0% purity, 7200 mesh), and C introduced as graphite (Cerac, 99.5% purity, \\\\37 7325 mesh). A series of samples in  \\\\37  the  system having  progressively  decreasing ZrC  contents were synthesized in order to determine the  e\\x80ect  of ZrC on  the  oxidation  behavior  of  the  materials. The ZrC/ZrB2 molar ceramics were approximately 10:1,  ratios  in the final  4:1,  1:1,  and  0:1, and the ZrB2/SiC ratio was materials and equal 2:1. The synthesis and densifi constant  for all  cation of ceramics were conducted during twostage hot pressing: at 1750(cid:14)C and 5 MPa for 1 h to and then at 2100(cid:14)C complete chemical reactions and 20 MPa for 0.5 h for ceramic densification.  All materials  were  characterized  by  density  (Archimedes method  and  helium pychnometry),  phase composition (X-ray di\\x80raction using a Sie mens  Theta-Theta Di\\x80ractometer),  and micro structure  (using  a  JEOL  JSM-6400V Scanning  Electron Microscope). For lattice parameter calcu lations in the HfC substoichiometric materials, an  internal silicon standard was used.  Mechanical  and  thermal  properties  evaluated  included elastic moduli,  tensile  strength, flexural  strength,  ductile-to-brittle  transition  temperature  (DBTT),  thermal  conductivity,  and  thermal  expansion  (CTE). The  test  specimens were  cut  using electrical discharge machining (EDM)  tech niques.  The thermal conductivity of the materials was measured (sample size: 1.91 cm 1(cid:2)1.27 cm thick)  using  a  comparative  rod  apparatus  technique,  which  compares  the  conductivity  of  a  specimen  with  the  conductivity  of  a  reference  (iron  and  stainless  steel  for  this  e\\x80ort). The  thermal  con ductivity was  calculated from the heat flow and  temperature di\\x80erence across a known gage length  within the specimen at a given thermal input. The CTE was measured (sample size: 0.64 cm 1(cid:2)5.08 cm long) up to 2500(cid:14)C using two di\\x80erent  dilatometers: a quartz-rod dilatometer provided data up to 1000(cid:14)C and a graphite dilatometer up to 2500(cid:14)C. The  same  specimens were used for both  temperature ranges.  The tensile test of OD(cid:2)4.06 cm ID(cid:2)1.27  ring  specimens  (4.57  cm  cm)  utilized  an  internal,  inflatable membrane to apply hydrostatic pressure.  Top  and  bottom enclosure  plates  and  annular  retaining rings were used to ensure that the loading  was applied in the radial direction.  Room and  elevated  temperature flexural tests 0.30(cid:2)0.635(cid:2)4.318  were  conducted  (sample  size:  cm)  to determine strengths, moduli, and DBTTs.  Flexural  testing was  conducted over a wide  tem perature  range  to determine  the  transition from  elastic stress-strain behavior  to elastic plus plastic  deformation. The DBTT was determined from the  change  in  slope  of  the  load-deformation  curve.  The elastic moduli were determined from the mea surement of ultrasonic  velocity,  as well  as  from  ring tensile and 3-point flexural tests.  The  oxidation  behavior  of  the ZrB2/ZrC/SiC characterized by thermogravimetric  ceramics was  analysis (TGA) using a TA Instruments SDT 2960  TGA/DTA module. The samples were heated at 20(cid:14)C/min up to 1500(cid:14)C in an argon/oxygen atmo sphere, which simulated the concentration of O2 in air. The TGA experiments were also conducted  isothermally for 5 h at a series of  temperatures.  The oxidation behavior of the ternary ceramics was  compared to that of pure ZrC and ZrB2 ceramics 2200(cid:14)C, prepared by hot pressing at 2300 and  respectively.  3 Results and Discussion  X-ray  di\\x80raction  (XRD)  of  the HfB2 showed that it contained 5% HfC, likely due to the  ceramics  excess of hafnium metal being  converted to the  carbide during hot-pressing. The HfN0(cid:1)92, HfC0(cid:1)98 and HfC0(cid:1)67 ceramics were all determined to be single phase. The measured lattice parameters of HfC0(cid:1)98 and HfC0(cid:1)67 were 4.64331 and 4.62368 A˚ respectively, correlating well with those reported in  ,  the literature.  All materials were well densified, as evidenced by  SEM of polished sections,  and density measure ments. An open porosity of <1% was measured  for all samples, and total porosity was determined  to be 5-10%. Figure 1 shows the microstructure of  the materials. The HfB2 ceramics [Fig. 1(a)] has a very fine-grained structure, with an average grain  size of 10-20 (cid:22)m, while the HfN0(cid:1)92 ceramics have a larger grain size of 40-60 (cid:22)m. Both materials  exhibit  similar  intergranular  fracture pattern. The  microstructure  of  the  hafnium carbide  ceramics  showed a substantial dependence on carbon stoi chiometry. The HfC0(cid:1)67 material sisted of very large crystallites (>200 (cid:22)m in size),  [Fig.  1(c)]  con while HfC0(cid:1)98  [Fig. 1(d)] has a smaller grain size  Table 1. The hot pressing parameters for processing billets  Material  Temperature ((cid:14)C)  Time (h)  Pressure (MPa)  HfB2 HfC0(cid:1)67 HfC0(cid:1)98 HfN0(cid:1)92  2160  3  27.3 16.5 31.7 34.1  2540  2 0.67  2500  2260  1  2406  M. M. Opeka et al.  \\x0c', '(40-60 (cid:22)m). The two carbide samples also exhib ited  distinctly  di\\x80erent  fracture  behavior.  The  HfC0(cid:1)67 while HfC0(cid:1)98 mode, as evidenced by the flatter fracture surface.  showed  extensive  intergranular  fracture,  failed in an transgranular  fracture  3.1 Thermal and mechanical testing results  The  thermal  conductivities of  the HfB2, HfC0(cid:1)98, HfC0(cid:1)67, and HfN0(cid:1)92 ceramics are shown in Fig. 2. The conductivity of HfB2 significantly exceeded that of the other materials: by a factor of 5 at room 800(cid:14)C. 2.5  temperature  and  by  a  factor  of  at  Additionally,  the  conductivity of HfB2 decreased with temperature, while that of the other materials  increased with temperature. Both carbides  exhib ited  the  same  temperature  coe(cid:129)cient  of  con ductivity,  but  the  conductivity  of  the HfC0(cid:1)98 the HfC0(cid:1)67 by a factor of two. This decrease in thermal conductivity can be due to  exceeded that of  increased phonon scattering from carbon vacancies  in the substoichiometric carbide.  The  thermal  expansion  curves  of  the  all  four  materials were very similar  in trend, as  shown in  Fig.  3.  The  expansion  of  all  materials was 1500(cid:14)C. Above  approximately 1500(cid:14)C,  the  same  up  to  the HfN0(cid:1)92 increase in the expansion rate.  ceramics  exhibited  some  Figure 4 shows the example of flexural stressdeflection curves for HfB2 at 1090 and 1230(cid:14)C. At 1090(cid:14)C, almost purely elastic behavior up to specimen failure was observed. In contrast, at 1230(cid:14)C,  the deflection includes a significant plastic defor mation component. Two modulus values were cal culated from this  curve based on the  elastic and  plastic  components. All modulus data calculated  from the flexural  testing are shown in Fig. 5. For  the HfC0(cid:1)98 observed up  sample, no plastic 2200(cid:14)C.  deformation was  to  In  contrast,  the  sub stoichiometric HfC0(cid:1)67 sample exhibited plastic low as 1090(cid:14)C. This di\\x80erence in deformation as  behavior can be attributed both to the significant  di\\x80erence  in microstructure  (Fig.  1)  and  to  the  compositional  di\\x80erence  (vacancy  concentration)  in  two materials. The HfB2 plastic deformation at a much lower  ceramics  exhibited  temperature  than the HfC0(cid:1)98 ceramics. The room temperature Youngs moduli  deter mined  by  all  three  techniques  for HfC0(cid:1)98, shown in Fig. 6. The tensile testing and  HfN0(cid:1)92 are based data  show very  good agreement with the  data gained from the ultrasonic velocity measure ments. The modulus data calculated from the flex ural  testing are lower (by approximately 15%)  for  all  four materials. The flexural  test-based moduli  Fig 1. Scanning electron micrographs of (a) HfB2, (b) HfN0(cid:1)92, (c) HfC0(cid:1)67, and (d) HfC0(cid:1)98.  Properties of refractory hafnium and zirconium compounds  2407  \\x0c', '2408  M. M. Opeka et al.  were expected to be low since the compliance of the  3.2 Oxidation Testing Results  load train is not accounted for  in the calculation  and  utilized  a  short-span  specimen. Of  all  four  materials, HfB2 HfC0(cid:1)67 the lowest moduli.  exhibited  the  highest  and  the  All ceramics in the ZrB2/ZrC/SiC system were fully densified and had predicted component ratios. The  microstructure of the ZrB2/ZrC/SiC ceramics (Fig. 7) is very fine-grained with bigger (5-10 (cid:22)m)  Fig. 2. The thermal conductivity of HfB2, HfN0(cid:1)92, HfC0(cid:1)67, and HfC0(cid:1)98.  Fig. 3. The CTE of HfB2, HfN0(cid:1)92, HfC0(cid:1)67, and HfC0(cid:1)98.  \\x0c', 'Properties of refractory hafnium and zirconium compounds  2409  crystals of ZrB2 crystals of SiC and ZrC. Compared to that,  surrounded by small  submicron  the  microstructure of  the pure ZrB2 ceramics is much courser with average crystal size of 30 (cid:22)m. The  presence of SiC and ZrC at  the grain boundaries  obviously inhibits excessive grain growth of ZrB2 during processing.  The oxidation behavior of  the ternary ceramics  was compared to that of pure ZrC and ZrB2. The pure ZrC (Fig. 8) was completely oxidized at 700(cid:14)C 20(cid:14)C/min.  TGA heating  continuous  during  at  Comparatively, the ZrB2 ceramics did not oxidize significantly even after 2 h heating at 1200(cid:14)C (Fig.  9). The  ceramics were protected by  liquid boria  Fig. 4. The flexural stress versus deflection of HfB2 at 1090 and 1230(cid:14)C.  Fig. 5. Young’s modulus versus temperature (from flexural testing) for HfB2, HfN0(cid:1)92, HfC0(cid:1)67, and HfC0(cid:1)98.  \\x0c', '2410  M. M. Opeka et al.  observed by SEM (Fig. 10) directly below the sur face between the ZrO2 grains, which is indicated by the meniscuses between the ZrO2 grains. Above 1200(cid:14)C, pure ZrB2 oxidized very actively due to the  intensive  evaporation of B2O3. However, tions showed the presence of about 10 wt% (15 vol%) of boria even at 1400(cid:14)C (Fig. 11), contra calcula dicting the literature data2,3 claiming that, at  tem Fig. 6. The room temperature moduli of HfB2, HfN0(cid:1)92, HfC0(cid:1)67, and HfC0(cid:1)98.  Fig. 7. SEM micrographs of ZrB2 and ZrB2/SiC/ZrC ceramics.  \\x0c', 'Properties of refractory hafnium and zirconium compounds  2411  1200(cid:14)C, peratures above the vapor pressure boria (m.p. 450(cid:14)C and b.p.(cid:25)1850(cid:14)C) becomes  of  so  significant that any boria formed during oxidation  of ZrB2 will completely evaporate. The SiC-containing ZrB2 ceramics (Figs. 8 and 9, curve with ZrC:ZrB2 ratio of 0:1) showed low and thermally stable oxidation rate up to 1500(cid:14)C. The  thickness of the oxidized layer in this material stays  1200-1400(cid:14)C after TGA heating for 5 hours (Fig.  12). For pure ZrB2, layer increases as a function of  the thickness of  the oxidation  temperature, con firming that boria alone does not prevent oxida tion.  The  stability  of  both mass  changes  and  thickness of oxidized layer with temperature  for  the ZrB2/SiC ceramics is an indication of very low rate of oxidation which is due to the formation of a  at about 100 (cid:22)m over  the  temperature  range of  protective layer of borosilicate glass on the surface  Fig. 8. Mass gain during TGA oxidation of ceramics in the ZrB2/SiC/ZrC system as a function of ZrC/ZrB2 ratio (ZrB2/SiC ratio is equal 2).  Fig. 9. Oxidation of ZrB2 and ZrB2/SiC ceramics in air (2 h).  \\x0c', '2412  M. M. Opeka et al.  of oxidized ceramics (Fig. 13). Compared to liquid  content. The low oxidation resistance of  the cera the borosilicate  B2O3 in ZrB2, containing ZrB2 higher melting temperature,  ceramics  glass  in the SiC has  higher  viscosity,  lower oxygen di\\x80usiv ity, and lower vapor pressure, thus providing much  mics having 10:1 and 4:1 ZrC:ZrB2 molar can probably be attributed to insu(cid:129)cient amounts  ratios  of ZrB2 and SiC to form a continuous protective surface layer of borosilicate glass during oxidation.  more e\\x80ective oxidation-protective capabilities.  The analysis of  the oxidation data for ZrC and  Practically similar oxidation rate of both ZrB2 and ZrB2/SiC ceramics below 1200(cid:14)C is attributed to the formation of the surface scale consisting of  ZrO2 and liquid B2O3 appreciably at these temperatures to participate in  since SiC does not oxidize  the formation of borosilicate glass. However, at temperatures above 1200(cid:14)C, the degree of SiC oxi led  ZrB2 describing the oxidation process. As  development  the  to  of  a  hypothesis  it was men tioned, pure ZrC was completely oxidized at 700(cid:14)C. The oxidation products are ZrO2 and carbon oxides. The ZrO2 forms a fine-grained porous scale which allows gaseous di\\x80usion of O2 through the pores to the ZrC surface, and therefore pro dation increases and becomes  su(cid:129)cient  to form a  vides no oxidation protection.  protective layer of borosilicate glass.  The  oxidation  materials  (Fig. 8)  stability  ZrB2/SiC/ZrC increased with increasing ZrB2  the  of  In case of ZrB2, the oxidation products are ZrO2 and liquid B2O3. In the initial stage of oxidation, boria fills all of the porosity and grain boundaries  Fig. 10. SEM micrographs of ZrB2 ceramics oxidized at 1300(cid:14)C for 5 h.  Fig. 11. Retained B2O3 content in the oxidation product of ZrB2.  \\x0c', 'Properties of refractory hafnium and zirconium compounds  2413  in  the  fine-grained  zirconia  forming  continuous  only 0.08 J/m2 compared to about 1 J/m2  for zir coating which  prevents  the  gaseous di\\x80usion  of  conia,  there is a large free energy driving force to  oxygen to the ZrB2 surface and provides oxidation protection. As boria evaporates readily at tem1100(cid:14)C,  condition  arises  in  peratures  above  a  keep the  surface of  zirconia  covered with boria,  providing an e\\x80ective oxidation protection barrier  even at higher temperatures.  which only a thin layer of boria covers  zirconia  The high surface free energy of  the fine-grained  particles. Since the surface free energy of boria is  zirconia explains the presence of boria in the oxi Fig. 12. Thickness of oxidized layer as a function of temperature after 5 h  isothermal heating.  Fig. 13. SEM micrographs of ZrB2/SiC oxidized at 1300(cid:14)C for 5 h.  \\x0c', 'dized layer  remaining constant at about 10 wt%  (15 vol % of the oxidation products) at 1200(cid:14)C. However,  tempera tures  above  regardless of  the  oxidation protection provided by liquid B2O3, ZrB2 oxidation process still occurs controlled by the di\\x80usion of oxygen through the boria  the  infil trated zirconia layer. The formation of borosilicate  glass in SiC-containing ZrB2 provides much more e\\x80ective oxidation-protective capabilities due to  higher viscosity, higher melting temperature,  lower  oxygen di\\x80usivity, and lower vapor pressure com pared to B2O3.  4 Summary  The  thermal  and mechanical  properties  of  haf nium-based nonoxide ceramics were measured.  It  was found that HfB2 had a much lower ductile-tobrittle transition temperature than HfN0(cid:1)92 or lowering the carbon  HfC0(cid:1)98. stoichiometry was also to decrease  The  e\\x80ect  of  the  transition  temperature. The  thermal  conductivity  of HfB2 than the  was much  greater  (by  a  factor  of  5)  carbides or nitride. The CTE of all materials tested were approximately the same up to 1500(cid:14)C,  with HfN0(cid:1)92 exhibiting a higher expansion than the others up to 2500(cid:14)C. The HfB2 ceramics had the highest modulus of the materials tested, while  HfC0(cid:1)67 measured during  had  the  lowest. While  the  modulus  the flexural  tests  tended to be  lower  than those measured by the ring and ultra sonic  methods,  the  trends  between  materials  was  consistent  regardless  of  the measurement  technique.  Ceramics in the ZrB2/ZrC/SiC system were prepared from mixtures of Zr (or ZrC), SiB4, and C using in-situ displacement reactions. The oxidation  behavior of  the  ceramics was  characterized as  a  function of phase composition which was varied by  changing the of component  ratios  in the reactant  mixtures. The SiC-containing ZrB2 ceramics had high oxidation resistance up to 1500(cid:14)C compared  to  pure ZrB2 and ZrC ceramics. The ZrB2/SiC ratio of about 2 (25 vol% SiC) is necessary for the  best oxidation protection. The presence of ZrC in  ZrB2 resistance.  ceramics  negatively  a\\x80ects  their  oxidation  A  hypothesis  describing  oxidation  behavior  of  the ZrB2/ZrC/SiC ceramics  is  pro posed.  Acknowledgements  The authors would like to thank Dr S. Dalleck and  Mrs K. Witkoski of NSWCCD for conducting the  TGA tests and assistance in SEM and EDS studies,  respectively.  The  authors  would  also  like  to  acknowledge N. Elsner  and D. Allen  of Hi-Z  Technology,  Inc. of San Diego, CA for  the pre paration of the hot-pressed billets. The NSWCCD  Independent Research program and Dr S. Fishman  of  the O(cid:129)ce of Naval Research provided support  for this project.  References  1. Fenter, J. R., Refractory diborides as engineering materi als. SAMPE Quarterly, 1971, 2(3), 1-15.  2. Tripp, W. C.  and Graham, H. C., Thermogravimetric  study of the oxidation of zirconium diboride in the temperature range 800(cid:14) to 1500(cid:14)C. J. Electrochem. Soc., Solid State Science, 1971, 118(7), 1195-1199.  3. Graham, H. C. et al., Microstructural  features of oxide  scales  formed on zirconium diboride materials.  In Cera mics  in Severe Environments, Proceedings  of  the  Sixth  University Conference on Ceramic Science, North Car olina State University  at Raleigh,  7-9 December  1970.  Materials Science Research, Vol. 5. Plenum Press, New  York, pp. 35-48.  2414  M. M. Opeka et al.  \\x0c']"
},{
  "_id": 125,
  "PDF": "Mechanism and Kinetics of Oxidation of ZrN Ceramics.pdf",
  "Text": "['Mechanism and Kinetics of Oxidation of ZrN Ceramics  Robert W. Harrison  †  and William Edward Lee  Department of Materials and Centre for Nuclear Engineering, Imperial College, London SW7 2AZ, United Kingdom  Oxidation of ZrN ceramics from 973-1373 K under static con ditions  reveals  parabolic rate behavior, indicative of a diffuIn-situ high temperature powder XRD  sion-controlled process.  found the oxidation mechanism begins with destabilization of  ZrN through formation of a ZrN1 - x phase with oxide peaks initially detected at around 773 K. The zirconium oxide layer was found to be monoclinic by in-situ XRD with no evidence of  tetragonal or cubic polymorphs present  to 1023 K. Bulk cera mic  samples  oxidized  at  1173  and  1273 K underwent  slower  oxidation than those oxidized at 973 and 1073 K. This change  in oxidation rate and hence mechanism was due  to formation  of a denser  c-ZrO2 polymorph stabilized by nitrogen defects. This N-doped dense ZrO2 layer acts as a diffusion barrier to oxygen diffusion. However,  at  an  oxidation  temperature  of  1373 K this layer is no longer protective due to increased diffu sion through it resulting in grain boundary oxidation.  I.  Introduction  Z IRCONIUM nitride is being considered for use in advanced nuclear power plants as inert matrix fuel or as accident tolerant fuel particle coatings due to its high thermal conductivity [45-50 W(mK)  \\x001],1,2  low neutron capture cross-section  and chemical compatibility with existing fuel ogy.3 Due to this  cycle  technol increased interest  in ZrN,  recent work has  focused on assessing thermophysical properties of ZrN and  mixed  phases  of  actinide  nitrides  dispersed  in ZrN.1,2,4-10  However,  there  have  been  few studies  on  the  oxidation  behavior of ZrN which is an important  factor  to be consid ered  especially  under  accident  conditions,  such  as  those  occurring in a loss of coolant accident  (LOCA).  The majority of oxidation studies of ZrN have been per formed on thin films due to the extensive use of ZrN as a hard coating on cutting tools.11-13 Krusin-Elbaum and Wittmer13 studied the oxidation kinetics of ZrN thin films (60 nm-1 lm) reporting the activation energy of oxidation to 241 \\x06 10 kJ/mol from 748-923 K with comprising monoclinic  be  the  oxide layer et al.12  and  cubic  ZrO2. showed the oxidation of ZrN coatings with thicknesses of 300 lm to have 773-1123 K and obey a parabolic rate law, in agreement with Krusin-Elbaum and Wittmer.13 Both studies show a pla Panjan  an  activation  energy  of  229 kJ/mol  from  teau region after  initial oxidation with both authors  report ing  parabolic  rate  behavior  and  it was  suggested  that  the  diffusion of oxygen through the oxide layer is the rate limitstudied the oxidation of 30 lm ZrN ing step. Caillet et al.14 coatings from 823-973 K observing linear kinetics suggesting  oxygen diffusion through the oxide layer does not affect  the  rate; however, a significant mass  loss was observed around  823 K. The authors propose nitrogen is  initially lost  leaving  hcp-Zr at the surface prior to ZrO2. Zirconium oxynitrides have been produced previously by mixing ZrO2 and ZrN powders15-17 nitriding ZrO2 under NH3 gas.18 Three oxynitride polymorphs have been et al.17 Zr2ON2, Zr7O8N4 and Zr7O11N2 reported by Gilles designated and , respectively. It can be  or  by  which  are  c,  b  b0  envisaged that  the  formation of an oxynitride phase would  be formed at the oxidation interface of ZrN, behaviour observed on oxidation of Si3N4.19,20 ZrC has similar chemical properties and crystal  similar  to  structure  to ZrN and its oxidation kinetics and mechanism in the bulk are better understood21-24 with onset of oxidation occurring between 653 and 750 K25 similar to that of ZrN.13 Oxidation  of ZrC above steps;21  743 K occurs  via  the  following  sequence of  1.  Formation of a ZrOxCy phase surface, which proceeds to form amorphous ZrO2 and free C. Crystallites of cubic zirconia (c-ZrO2) layer with free  2.  nucleate  and  grow to form a dense oxide  carbon,  which stabilizes the c-ZrO2. Oxygen then diffuses through this oxide layer to the free  3.  carbon producing CO2 which escapes via any existing cracks or pores leaving behind voids and pores in the  zirconia layer which, with less carbon left  to stabilize it  transforms to monoclinic zirconia (m-ZrO2) with small amounts of tetragonal polymorph (t-ZrO2).25  ZrN was expected to show similar oxidation properties as  ZrC and the aims of  this work were to examine the mecha nism of oxidation of bulk ZrN samples, determining phases  present  in the oxidation layer and investigate the existence of  different  layers at  the  reaction interface,  such as oxynitride  or high temperature ZrO2 polymorphs as observed with ZrC ceramics.25 Oxidation temperatures were chosen as to be able  to examine the initial to intermediate stages of oxidation. The higher temperatures (1073-1373 K) are also in the range  of the expected nominal operating temperature and initial temperature increases in the core following a LOCA.26 Static  atmosphere was used to aid examination of  the  initial and  intermediate steps of oxidation and also as  this  is envisaged  as being more akin to a LOCA whereby the coolant  system  fails but  the nuclear reactor core remains intact.  II.  Experimental Procedure  Commercially-available ZrN powder (1-2 lm, ≥99.0%; Sigma  Aldrich, Gillingham, UK) was densified using spark plasma  sintering (HP D/25/1; FCT systeme GmbH, Rauenstein, Ger many)  in  a  graphite  die  (30mm diameter),  under  vacuum  (0.5 mbar). Sintering was carried out for 10 min at 2373 K (+100 K/min) under 50 MPa uniaxial pressure. The resulting 7.13 \\x06 0.01 g/cm3 calculated from the average of three measurements performed  bulk  ceramic  samples  had  a  density  of  by 97.7 \\x06 0.1%. the Archimedes method In-situ phase  and  a  theoretical  density  of  analysis of ZrN powder  (1-2 lm, ≥99.0%; Sigma Aldrich) under static air was done using high tempera ture  (up  to  1023 K) X-ray  diffraction  (HT-XRD; X’pert  N. Jacobson—contributing editor  Manuscript No. 36116. Received November 4, 2014; approved February 24, 2015.  †  Author  to whom correspondence  should  be  addressed.  e-mail:  r.harrison11@  imperial.ac.uk  2205  J. Am. Ceram. Soc., 98 [7] 2205-2213 (2015)  DOI: 10.1111/jace.13575  © 2015 The American Ceramic Society  Journal  \\x0c', 'MD,  Phillips, Amsterdam,  the Netherlands)  in  a  furnace  using a Pt heating strip (HDK 2.4; Buhler, Tubingen, Germany) at 40 kV and 40 mA,using Ni-filtered CuKa radiation between room temperature and 1023 K, with scans  performed every 12 K between 673 and 1023 K. Due to the  nature of  the  in-situ XRD requiring  thin layers of  sample  experiments could not be done on bulk samples and so were  done on to powders. Ceramic samples with dimensions 3 mm 9 3 mm 9 3 mm  hand abraded with 600 grit SiC paper were oxidized under  static air using thermogravimetric analysis  coupled differen tial  thermal analysis  (TGA/DTA)  to measure  the  change  in  weight as a function of  time  (STA-449-F1; Netzsch GmbH,  Selb, Germany) at 5 temperatures; 973, 1073 and 1173, 1273 and 1373 K (+20 K/min)  for a dwell  time of 5 h. Change in  mass was  continually  recorded  using  the micro-balance  of  the TGA, weight  change  after  0,  1,  2,  3,  4,  and 5 h time  points  is  reported. Oxidation using  the TGA ensures more  reliable results in weight change as any oxide scale spallation  remains  in the  crucible  and there  is no excess weight  gain  from the  cooling of  the  furnace at  the desired time points.  Experiments were performed once at each isothermal oxida tion  temperature, so that results are directly of MW/A, (MW/A)2 (MW/A)3 made to determine which equation gave the best according to Eqs. 1, 2, and 3 (where MW is the change in weight (mg) A is the area of the pellet (cm2), kl, kp and kc are the linear, parabolic and cubic rate constants respectively is the time (s)).20,27,28 Errors of MW/A are given as the percentage error of the TGA/DTA instrument compounded  comparable.  Plots  and  against  time were  linear fit  and t  with the error from the area measurements of  the pellets.  ðMW  A  Þ ¼ kl t  (1)  ðMW  A  Þ2 ¼ kp t  (2)  ðMW  A  Þ3 ¼ kc t  (3)  Microstructural  characterization of  the oxidized, polished  pellets was performed using secondary electron imaging (SEI)  in an SEM (JSM-6400; JEOL, Tokyo, Japan) equipped with  energy dispersive  spectroscopy  (EDS) detector  for  chemical  analysis  (ultra thin polymer window,  INCA; Oxford instru ments, Oxford, UK). Focused ion beam (Helios Nanolab 600;  FEI, Hillsboro, OR) was used to prepare samples and trans mission electron microscope  (TEM) analysis was performed  on the oxidized ceramics (HRTEM FX2000, JEOL). Selected  area diffraction (SAD) patterns were indexed by matching dhkl values with reference patterns and calculated SAD patterns  using  SingleCrystal  software  (version  2.2.9  CrystalMaker  Software Limited, Oxfordshire, UK). Lattice parameters were  calculated  from the  indexed  SAD patterns  along with  the  measured dhkl spacing’s using Eq. (4) and contain a 2% error due to the small (0.5-50) angle of diffraction.29  a ¼ d  p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi  h2 þ k2 þ l2  (4)  III.  Results and Discussion  (1)  Oxidation Kinetics  In-situ XRD of ZrN powder during oxidation is presented in  Fig. 1. Monoclinic ZrO2 no dwell time and was the only polymorph detected during  is  initially detected at 773 K after  the oxidation process which as expected since the tetragonal and cubic forms are both high temperature (≥1450 K) poly morphs. However,  this is not  in agreement with observations  of a very low intensity reflections identified as a minor cubic  ZrO2 polymorph detected by glancing angle XRD of dized thin films by Krusin-Elbaum and Wittmer.13  oxi It  should  be noted that  the observation of  low amounts of cubic ZrO2 would be difficult by in-situ XRD due to the differences in  experimental  technique with glancing  angle diffractometers.  Peak intensity is also lower in in-situ XRD compared to typi cal diffractometers due  to the  scattering  from the  chamber  window of  the furnace between the source and detector. Fig ure  2  shows an excerpt of the (200) ZrN reflection which initially the peak shifts to lower 2h angles, corre reveals that  sponding to an increase  in lattice parameter due  to thermal  expansion until around 673 K. Above 673 K the peaks begin to higher 2h angles  to shift  suggesting a decrease  in lattice  parameter which could arise from loss of nitrogen from the lattice resulting in formation of ZrN1\\x00x. Further temperature leads to the peak splitting, resulting in a peak  increase of  consisting of  two main components  (observed for  all ZrN  reflections). This suggests the existence of two or more ZrN phases, such as a substoichiometric ZrN1\\x00x phase and Zr-O- N phase. Incorporation of oxygen into the ZrN lattice to  form an  oxynitride with  the  same  crystal  structure would  result  in the  formation of a  smaller  lattice  resulting  in the  Fig. 1. In-situ XRD characterization of ZrN powder from RT1023 K with scans performed every 100 K between RT-1023 K with  no dwell with  no dwell times at with monoclinic ZrO2 40 and ZrN31.  the  temperatures. Peaks  indexed  Fig. 2.  Excerpt of  in-situ XRD of 220 ZrN reflection from RT 1023 K, scans performed every 100 K between RT-673 K and every 12 K between 673-841 K with no dwell time.  2206  Journal of  the American Ceramic Society—Harrison et al.  Vol. 98, No. 7  \\x0c', 'to higher 2h. This agrees with the observations of peak shift Gendre et al.30 who showed the lattice size of ZrC decreased  with  oxygen  content  to  a  stoichiometry  of  around  ZrC0.79O0.13.  The  change  in weight normalized for  the  surface area of  the ceramic samples as a function of  time is shown in Fig. 3  which reveals a non-linear relationship [Eq. (1)] for all (indicated by R2 values ≤0.99 for a linear fit except for the sample oxidized at 1273 K which fits well  sam ples  in Fig. 3),  to a  linear  relationship. However,  this may be because the oxida tion rate  is  still  in the  linear  region  and  longer  oxidation  times may  show parabolic  rate behaviour. Figure  4  shows  the square of  the weight gain normalized for  surface area as  a function of time for each isothermal oxidation temperature with all R2 values ≥0.99 showing a good fit. All samples had a R2 ≤0.97 for a cubic relationship [Eq. (3)] showing a poor  fit. A parabolic  rate  is  expected  for  the  samples  due  to  increased diffusion distance for oxygen through the growing  oxide layer to the ZrO2/ZrN reaction interface. It can be seen from Fig. 3 and plotting the rate constant for each tempera ture (Fig. 5)  that  the samples oxidized for 5 h at  lower  tem peratures  (973 and 1073 K)  showed more weight gain than  those oxidized for  5 h at  1173  and 1273 K. The  rate  con stants  in Fig.  5  show a  rapid  increase  between  973  and  1073 K followed by a large decrease at 1173 K and then no  significant difference between samples oxidized at 1173 and  1273 K. At  the higher  temperature of 1373 K however,  the  rate of oxidation is  the  fastest observed for  all  isothermal  oxidation temperatures. The proposed change in mechanism  between  low and  high  temperature  oxidation will  be  dis cussed further in Section III(3).  (2)  Oxidation Layer Characterization  On the ceramic sample oxidized for 5 h at 973 K (Fig. 6) an  outer layer of brittle, porous ZrO2 was formed which spalled away from the surface of the sample, revealed by loose pow der  present when  removing  the  sample  from the  crucible.  Figure 6(a) also shows  light  regions with increased porosity  further  into  the  bulk  from the  surface  that EDS  revealed  contained oxygen and nitrogen,  indicating the oxidation layer  penetrated further than the surface layer. This is more clearly  observed in the  sample oxidized at 1073 K for 5 h (Fig. 7)  where a dense but cracked layer of ZrO2 at sample is observed on top of a cracked  the surface of  the  layer with mixed  porosity which EDS also showed contained oxygen. SEIs of  the samples heated to 1173 and 1273 K for 5 h are shown in  Figs. 8 and 9, respectively. Figure 8 reveals three layers, EDS  of  the first  two revealed only Zr and O, while the third was  unreacted bulk ZrN. The presence of carbon arises  from the  epoxy resin used to mount  the  samples. The SEIs  in Fig. 9  reveal  that  the  grains  closer  to  the  surface  (nearer  to  the  oxide  layer) are  readily  visible; however, grains  toward the  bulk are not. Line  scans of Fig.  9(b)  reveal  the Zr  counts  decrease in the crack in the outer oxide layer, however,  they  do not decrease significantly at  the revealed grain boundary.  This  suggests  the grains have been revealed by oxygen diffu sion  along  the  grain  boundaries  nearer  to  the surface. A 20 lm thinner) and 1173 K for 5 h  decrease  in  oxide  layer  thickness  (around  between the  samples oxidized at  1073  was observed. The oxide scale thickness did not  show a sig nificant  increase  for  the  samples  oxidized  at  1173  and  1273 K for 5 h. SEIs  (Fig. 10)  reveal  the sample oxidized at  1373 K for 5 h showed extensive grain boundary oxidation  in the bulk. The oxide scale is also cracked and spalled away  from the sample on handling,  similar  to the sample oxidized  at 973 K for 5 h.  Transmission electron microscope of a FIB section of  the  intermediate layer observed in the sample oxidized at 1173 K  is  shown in Fig. 11(a). The dark phase  is a thick region of  monoclinic ZrO2 as characterized by XRD (Fig. 1) and corresponds to the surface of the oxidized sample. The bulk of  the  sample  is ZrN and the SAD pattern from this  region  matched ZrN [Fig. 11(b) which indexed as the [001]ZrN zone  axis]. 4.400 \\x06 0.090 \\x17A is However,  the  calculated  lattice  parameter  of  smaller  than the literature value given at  Fig.  3.  Change  in mass  per  unit  surface  area  of  ceramic ZrN  samples as a function of time for isothermal oxidation experiments.  Fig. 4.  Square of change in mass per unit  surface area of ceramic  ZrN samples  as  a  function  of  time  for  isothermal  oxidation  experiments.  Fig. 5.  Parabolic  rate  constants  (kp) as a function of for ceramic ZrN samples oxidized for 5 h.  temperature  July 2015  Oxidation of ZrN  2207  \\x0c', '2208  Journal of  the American Ceramic Society—Harrison et al.  Vol. 98, No. 7  (a)  (b)  Fig. 6.  layer.  SEIs of oxide layer of ZrN sample heated to 973 K for 5 h,  (a) shows fragmented oxide layer,  (b) shows grains revealed near to oxide  (a)  (b)  Fig. 7.  SEIs of oxide layer of ZrN sample heated to 1073 K for 5 h, (a) shows cracks present in bulk ZrN, (b) shows cracks in oxide layer.  4.574 \\x17A.31 The decrease in lattice size may be due to nitrogen vacancies and oxygen content causing the lattice to shrink.  SAD patterns  of  the  intermediate  grain  between  the  bulk  ZrN grain and the surface ZrO2 are shown in Figs. 11(c) and  (d), which has a width of around 1 lm and agrees well with  the width of  the  layer  in Fig. 8. The SAD pattern in Fig.  11(d) indexed as (\\x16311)/(002)  the  the  c-ZrO2 d-spacings (1.678) matches with  zone axis as  [103]  the  ratio of  the  cubic  \\x0c', 'July 2015  (a)  Oxidation of ZrN  (b)  2209  Fig. 8.  SEIs of oxide layer of ZrN sample heated to 1173 K for 5 h,  (a)  shows  region where intermediate layer  is penetrated by oxide scale,  (b) shows cracks in the oxide layer.  (a)  (b)  Fig. 9.  SEIs of oxide layer of ZrN sample heated to 1273 K for 5 h,  (a) shows distance into the bulk that grains have been revealed,  (b) shows  cracks in oxide scale, EDS line scans show zirconium counts (bottom left) and oxygen counts (bottom right).  pattern (1.658) as opposed to the ratio of 1.728 in the tetrag [103] zone axis. The lattice parameters calculated  onal ZrO2 4.877 \\x06 0.100 from the SAD patterns 4.877 \\x06 0.100 \\x17A in Figs. although \\x19 0.2 \\x17A smaller than the cubic ZrO2 lattice paramand respectively, which 5.090 \\x17A32 match eter of better than with the ZrN lattice  11(c)  and  (d)  were  parameter of 4.574 \\x17A.31 Figure 12(a) shows a TEM image of a FIB section a grain boundary from the sample oxidized at  1373 K for  5 h, which  reveals  the material  between  the  ZrN grains  in Fig. 10 to consist of  sub-micron grains,  inter granular  cracking  can  also  be  seen. The  SAD pattern  of  the sub-micron grains  shown in Fig. 12 indexed as  the cubic  \\x0c', '2210  Journal of  the American Ceramic Society—Harrison et al.  Vol. 98, No. 7  (a)  (b)  Fig. 10.  SEIs of oxide layer of ZrN sample heated to 1373 K for 5 h,  (a) shows oxide layer and bulk ZrN grains that have undergone attack  by oxidation, (b) higher magnification of ZrN grain revealed by EDS to have oxygen content between ZrN grains.  (a)  (c)  (b)  (d)  Fig. 11.  (a) BF-TEM image of FIB section taken from intermediate  layer of  sample oxidized at 1173 K for 5 h,  (b) SAD pattern of bulk  region, c and d) SAD patterns of intermediate layer grains.  [1\\x1611] 4.850 \\x06 0.100 \\x17A. Although nitrogen has a larger atomic radii ZrO2 zone axis and had a lattice parameter of than oxygen due to its weaker effective pull on its outer elec trons  a decrease  may be due  in observed  cubic ZrO2 to oxygen vacancies formed by substitution of  parameter  lattice  oxygen  with  nitrogen.  Evidence  for  ordering  of  oxygen  \\x0c', 'vacancies and nitrogen impurities can also be observed from  the forbidden reflections present  in Figs. 11(c) and (d) and 12  which have d-spacing values which are half of  the fundamen tal  refection  spots,  suggesting  they  are  from superlattice  reflections and not  from neighboring crystals of differing lat tice  size.  Similar  forbidden  reflections  and  stabilization  of  c-ZrO2 by the addition of ZrN were observed by Van Tendeloo and Thomas.15 The authors state that the vacancies  introduced  by  replacing  oxygen with  nitrogen  order  along  one of  the  threefold axes which lowers  the  symmetry of  a  rhombohedral Zr7O11N2 phase which is closely related to the cubic fluorite structure allowing the forbidden reflections. It  is plausible a similar phase may be observed in this work,  however, only the observation of  stabilization of c-ZrO2 was  made with certainty.  (3)  Discussion of Oxidation Mechanism  Parabolic  rate behavior was observed for all oxidation tem peratures  except  the  sample  oxidized  at  1273 K for  5 h,  indicative  of  a  diffusion  controlled  rate  limiting  step. The  observed  decrease  in  rate  of  oxidation  between  1073  and  1173 K along with the difference in morphology of  the oxi dation layer  cross-sections  (e.g. Fig. 7 compared to Fig. 8)  and the presence of a high temperature c-ZrO2 phase at oxidation layer (Figs. 8, 11, and 12) suggest the mechanism  the  of oxidation changes with  temperature  and  a  schematic  is  shown in Fig. 13. At  temperatures up to 1073 K a porous  oxide  layer  forms, which would readily allow oxygen diffu sion through pores and voids and nitrogen (or NOx species) diffusion out of the sample. The increase in oxidation rate  (Fig. 5)  is  consistent with the porous microstructure of  the  oxide  layers  (Figs.  6  and 7). At oxidation temperatures of  1173  and  1273 K a  dense  intermediate  oxide  layer  forms  which inhibits  the diffusion of oxygen and therefore the rate  of oxidation. This dense  layer  is  revealed to be  c-ZrO2 by 11). As the  TEM analysis  of  the  intermediate  layer  (Fig.  sample  cools  the oxide  layer  transforms  to the monoclinic  polymorph  creating  the  cracks  present  in  the  dense ZrO2 the surface (Figs. 8 and 9) due to the volume  polymorph at  change  associated  with  the  polymorphic  transformation.  However,  the  sample oxidized at 1373 K for 5 h showed a  rapid increase  in rate of oxidation.  It  can be  seen from the  SEIs  (Fig. 10)  that  there has been extensive grain boundary  oxidation which would result  from the increased diffusion of  oxygen through the oxide scale. Shimada and Ishil21 describe a change  in the mechanism  of ZrC oxidation with increasing temperature. At higher temperatures (≥743 K), submicron c-ZrO2 crystals nucleate from the amorphous ZrO2 + C layer and this forms a compact  (a)  (b)  Fig.  12.  (a) BF-TEM image of FIB section taken from grain boundary of  sample oxidized at  1373 K for  5 h,  (b) SAD pattern of  area  indicated with circle.  Fig. 13.  Proposed mechanisms of oxidation of ZrN at low (≤1073 K) and higher (≥1173 K) temperatures.  Table I .  Molar Volumes and Thermal Expansion  Coefficients of ZrN, 3 mol% Yttria-Stabilized ZrO2, and m-ZrO2  40-42  Phase  Molar volume (Vm) \\x001) (cm3 mol  Coefficient of  thermal expansion \\x006K \\x001)  at 1273 K (10  ZrN  14.4  7.6  c-ZrO2 m-ZrO2 3 mol% Y2O3-ZrO2  19.6  — —  21.0  —  10.8  July 2015  Oxidation of ZrN  2211  \\x0c', 'dense  layer whose  formation becomes  the  rate determining  step for oxygen diffusion. However, at higher  temperatures  (823 K)  and  longer  oxidation  times,  the  small  crystallites  grow and the volume expansion causes intergranular fracture  providing paths  for oxygen diffusion in and CO2 diffusion volumes of ZrN, m-ZrO2 and c-ZrO2 are presented in Table I along with the thermal expansion data for  out. Molar  ZrN and  yttria-stabilized  ZrO2. is assumed  Thermal  expansion  data  for  ytrria  stabilized ZrO2 stabilized ZrO2 polymorph in this work. From the values Table I, it can be seen the associated volume change between  to  be  similar  to  the  in  phases and the  thermal  expansion mismatch between stabi lized ZrO2 could both provide sources of stresses and result in intergranualr fracture. Puclin and Kaczmarek33  examined  the growth in crystal  size of ZrO2 t-ZrO2 1223 K and m-ZrO2 can also be envisaged as a source of stress and cracking due  reporting a rapid increase  in size of  at  at  1273 K which  to the growth of  submicron size of  the c-ZrO2 grains  in the  grain boundary. From the observation of a thin (≤1lm) c-ZrO2 layer in the (≤1lm) ZrN samples in this work (Fig. 8) and small grains observed by TEM (Fig. 11) a similar rate limiting step can be  envisaged for ZrN, which causes decreases in the rate of oxida tion between the samples oxidized at 1073 and 1173 K. How ever,  these  crystallites  then  grow and  could  transform to  m-ZrO2 causing cracking between the grains, as observed in Fig. 12(a), where it can be seen the crack follows what was the  ZrN grain boundary in the oxide layer and slightly into the bulk. Berkowitz-Mattuck34 also observed that oxidation of temperature ≤1470 K resulted in porous m ZrC samples at  ZrO2 which fractured due to intergranular oxidation; however, above 1470 K a denser ZrO2 layer formed slowing diffusion of oxygen to the bulk attributed to sintering effects on the oxide  layer. The observations of denser outer layers of ZrO2 of samples oxidized 1173 and 1273 K (Figs. 8 and 9) agree well with the dense ZrO2 layer reported by Shimada and Ishil21 and Berkowitz-Mattuck.34 Simultaneously, the formation of dense  the  scales of ZrO2 has been proposed to provide effective oxidation resistance in ZrB2-SiC35 and ZrC-SiC36 composites. Different mechanisms are thus proposed for the low temperature (≤1073 K) and higher temperature (≥1173 K) oxidaoxidation mechanism ≤1073 K can  tion  of ZrN. The  be  described in several  steps,  initially oxygen diffusion through  the surface ZrN grain boundaries  results  in a ZrOxNy phase which then forms m-ZrO2. The formation of this oxide layer results in a diffusion-controlled process as observed by  the  parabolic kinetics shown in Fig. 4. Stresses from the m-ZrO2 grains between the ZrN grains result in intergranular fracture  and the creation of a porous oxide layer as observed in Fig.  6, similar to the observation of Berkowitz-Mattuck ZrC.34 Figure 7 also reveals lighter regions in the bulk of  for  the  ZrN which  correspond  to  oxygen-rich  areas  formed  from  grain boundary diffusion of oxygen that have not yet caused  fracture.  The  creation  of  these  grain  boundary  fractures  opens channels  for  further  ingress of oxygen and removal of  nitrogen advancing the rate of oxidation, matically in Fig. 13 labeled ≤1073 K.  this is  shown sche Oxidation  of  ZrN  at  higher  temperatures,  however,  proceeds  by  the  formation  of  a  high  temperature  c-ZrO2 This high  polymorph  comprised  of  submicron  grains.  temperature phase is stabilized by the substitution of oxygen  with nitrogen-forming oxygen of ZrN is exothermic,37 it  vacancies. As  the  oxidation  is possible the local  temperature at  the  interface will be higher  than the oxidation temperature.  This increase in temperature may give rise to sintering of  the  c-ZrO2 crystals similar to effects observed in the oxidation of ZrC by Berkowitz-Mattuck34 however, the nominal sintering temperature of ZrO2 is around 1673 K.38,39 This dense cZrO2 is also similar to that reported by Shimada and Ishil.21 The dense oxide layer inhibits oxygen’s ingress and therefore  reduces  the  rate of oxidation, providing  a diffusion barrier  resulting  in  parabolic  behavior  and  a  lower  oxidation  as  shown by the rate constants in Fig. 5. However,  from Fig. 9,  grains are more readily observed suggesting greater potential  for oxygen diffusion via grain boundaries. Expansion of  the  c-ZrO2 crystallites or gen defects are removed results  transformation to monoclinic as nitro in the  cracks present  in the  oxide layers in Figs. 8 and 9. The sample oxidized at 1373 K  for  5 h  showed  an  increase  in  oxidation  rate  and  this  is  attributed  to failure  of  the dense oxide  scale  as,  although  oxygen  diffusion  through  the  dense  zirconia  layer will  be  slow,  it will  increase with temperature. This  is observed as  the increase in grain boundary oxidation of samples oxidized  at  1273  and 1373 K for  5 h (Figs.  9  and 10  respectively).  Thus,  eventually  the protective oxide  scales  in the  samples  oxidized at 1173 and 1273 K will  fail with time as oxygen  diffuses  further  through the dense oxide layer.  Intergranular  cracking observed in the TEM image of  the sample oxidized  at 1373 k for 5 h (Fig. 12) will  create  channels  for  further  oxygen diffusion into the sample.  IV.  Conclusions  The oxidation kinetics and mechanism of bulk ZrN ceramics  between 973 and 1373 K have been studied.  In-situ powder  XRD reveals oxidation begins by  the  removal of nitrogen  from the ZrN lattice (forming two or more phases with smal ler  lattice  size  indicating  either  substoichiometric ZrN or  a  ZrN phase with oxygen defects) and proceeds by the forma tion  of  a  porous m-ZrO2 1023 K. Oxidation rates in bulk ceramics decreased at  phase  at  temperatures  up  to  tem peratures at or above 1173 K attributed to formation of a  higher  energy c-ZrO2 phase which is proposed to slow oxygen diffusion by forming a dense layer, limiting the diffusion  of oxygen to the bulk. At 1373 K however,  this barrier  to  oxygen  diffusion  fails  due  to  increased  oxidation  of  grain  boundaries  and intergranular  cracking  creating more  chan nels  for oxygen diffusion. The nature of  the protective scale  observed for  ceramic  samples oxidized at 1173 and 1273 K  for 5 h needs  further  investigation,  examining longer oxida tion times and oxygen pressures.  Acknowledgments  The authors are grateful  to the EPSRC for their financial support of  this pro ject  (grant EP/J500239/1). We also thank Dr D. Horlait and Dr. S. Humphry Baker  for  their  invaluable discussions and Mr Richard Sweeney for his help  with the HT-XRD measurements.  References  1R. Harrison, O. Ridd, D. D. Jayaseelan, and W. E. Lee, “Thermophysical  Characterisation of Ceramics Fabricated via Carbothermic Reduction-Nitridation,” J. Nucl. Mater., 454 [1-3] 46-53 (2014). 2M. Pukari, M. Takano, and T. Nishi, “Sintering and Characterization of (Pu,Zr)N,” J. Nucl. Mater., 444 [1-3] 421-7(2014). 3M. Burghartz, G. Ledergerber, H. Hein, R. R. van der Laan, and R. J. M.  Koningss, “Some Aspects of the Use of ZrN as an Inert Matrix for Actinide Fuels,” J. Nucl. Mater., 288 [23] 233-6(2001). 4A. Ciriello, et al., “Thermophysical Characterization of ZrN and (Zr,Pu)N,” J. Alloys Compounds, 473 [1-2] 265-71 (2009). 5Y. Arai, M. Akabori, and K. 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Sglavo, “Electric Field Assisted Sintering of Cubic C,” J. Am. Ceram. Soc., 96 [5] 1342-4 (2013). Zirconia at 390 40J. D. Hanawalt, H. W. Rinn, and L. K. Frevel, “Chemical Analysis by X-ray Diffraction,” Indus. Eng. Chem. Anal. Ed., 10 [9] 457-512 (1938). 41M. Takano, M. Akabori, Y. Arai, and K. Minato, “Thermal Expansion  in SOFC Manufacturing,”.  ∘  of TRU Nitride Solid Solutions as Fuel Materials for Transmutation of Minor Actinides,” J. Nucl. Mater., 389 [1] 89-92,  (2009). Thermochemistry and Ther mophysics of Nuclear Materials Proceedings of  the Twelfth Symposium on  Thermochemistry and Thermophysics of Nuclear Materials. 42H. Hayashi,  for Various Yttria Contents,” Solid State  “Thermal Expansion Coefficient of Yttria Stabilized [5-6] 613-9  Ionics,  176  al.,  et  h  \\x0c']"
},{
  "_id": 126,
  "PDF": "Mechanisms for Variability of ZrB2‐30 vol_ SiC Oxidation Kinetics.pdf",
  "Text": "['Mechanisms for Variability of ZrB2-30 vol% SiC Oxidation Kinetics  †,* Brandon Patterson, David Lichtman, Siying Liu, and Elizabeth Opila** Kathleen Shugart,  Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22904  The  oxidation  kinetics  of ZrB2-30 vol% SiC were goal of understanding the underlying  analyzed  statistically  with  the  mechanisms  for observed variability. A box furnace was used  to oxidize specimens for times between 30 s and 100 h at temperatures of 1300°C-1550°C in air. The specimens were char acterized  to  determine weight  change,  scale  thickness,  and  scale  composition to quantify the oxidation behavior. Weight  gain measurements of different specimens after showed differences of up to 2 mg/cm2  100 min  of  exposure  for  the  same  testing conditions where the average weight gain was 2.54 mg/ 30%-80% was cm2. Variation  of  observed  in  the  average  thickness of each layer of  the oxide within a single  specimen.  Viscous  glass flow was  ruled  out  as  a  potential mechanism.  Glass  bubble  formation was  proposed  as  the main  cause  for  oxidation kinetics variability.  I.  Introduction  U LTRAHIGH-TEMPERATURE Ceramics  (UHTCs) are a fam ily  of materials which 3000°C, making  have melting  temperatures  in  excess  of  them attractive  for  aerospace  applications  such as Thermal Protection Systems  (TPS) on  hypersonic  flight  vehicles. Future  hypersonic  vehicles will  likely  have  sharper wing  leading  edges  and  sharper noses, speeds.1,2  allowing  for  better maneuverability  and  higher  When the  radius of  the wing  tip is made  sharper  and the  speed is increased the temperatures seen by the vehicle will increase to 1700°C-3000°C due to shockwaves being in closer surface.2-4 beyond the operating temperatures of current TPS, necessitating the development of new TPS materials.3-6 One of largest concerns for nonoxide TPS materials is oxidation  proximity  to  the  vehicle  Such  temperatures  are  the  behavior at ultrahigh temperatures. Current materials, including SiC-coated carbon-carbon composites, do not have  sufficient oxidation resistance at these ultrahigh temperatures liquid (SiO2 Tm = 1723°C) and gaseous due to formation of [SiO(g)] products.4 UHTCs, on the other hand, form solid  oxides which  themselves  have  very  high melting  tempera tures. One member of  this family, ZrB2, has been extensively studied since the 1960s due to its high melting temperature (Tm = 3245°C), cm3), strength retention at high temperatures, chemical bility, and resistance to erosion/corrosion.3,4,7,8 However,  relatively  low  theoretical  density  (6.09 g/  sta the  oxidation rate of ZrB2 is rapid and unacceptable for reusable applications. Additions of SiC have been shown to improve the oxidation behavior of ZrB2.9-11 The oxidation/vaporization of ZrB2-SiC follows tions:  the reac ZrB2 (s) þ 5 2  O2 (g) ¼ ZrO2 (s) þ B2O3 (l)  (1)  SiC(s) þ 3 2  O2 (g) ¼ SiO2 (s) þ CO(g)  (2)  B2O3 (l) ¼ B2O3 (g)  (3)  When exposed to temperatures  above  1300°C, ZrB2-SiC grains above the base  forms  a  two layer oxide, with ZrO2 material and borosilicate glass filling in the  spaces between layer.12,13 CO  the grains and providing a protective  surface  (g) diffuses outward through the  scale; however, accumula tion of large pressures glass.14 B2O3(l) also has a high B2O3 in the borosilicate glass.15 Most studies of ZrB2-SiC oxidation are material and repeat tests using any one set of conditions  can  cause  bubbles  to  form in  the  vapor  pressure,  depleting  screen ings  and  material are  rarely reported.  It has been noted that while a  general  behavior  is  common  for  certain  time/temperature  conditions, quantitative agreement between studies is not necessarily seen.13  for  even  nominally pure diborides  It  is also  difficult  to  compare  oxidation  results  from one  study  to  another due  to changing  and often unspecified parameters.  Parameters  that  are  often  not  reported  include  the  atmo sphere under which the material  is tested, gas flow rate, spec imen  texturing,  and  humidity.  Some  experiments  are  conducted isothermally while others temperature.16-18 apparatuses are used, which can impact  ramp up to the desired  In addition,  a number of different  testing  the  results. These  oxidation  methods  include  box  furnaces,  oxyacetylene  torches,  arc-jet  facilities,  and  resistive  heating.12,17,19,20  Finally, results are often reported using different conventions,  such as weight gain, scale thickness, or recession as functions time and temperature.13  of  While  several models  for UHTC oxidation kinetics based  on results  found in the  literature  exist,5,21,22  the wide  range  of results makes accurate modeling difficult. The model by Parthasarathy et al.13,23  suggests  the quantitative differences  from different  tests could be due to small changes in porosity  or trace impurities in the specimens, both of which are often unreported. Levine et al.18 suggest  that variation in mechani cal properties may be due to issues  in repeatable processing,  and this could also lead to differences  in oxidation behavior.  Other sources of variability in oxidation kinetics may include  poor experimental or manufacturing technique, differences in  relative  humidity  during  testing,  gas  flow during  testing,  impurities  (either  environmental  or  processing),  surface  roughness, viscous flow on the macroscale, or  the formation  of bubbles in the glass scale.  This work includes a statistical analysis of ZrB2-30 vol% SiC oxidation with reporting of all testing conditions. Varia tion in oxidation results are examined in relation to viscous  flow and bubble  formation in the glass  scale. A companion  paper  examines the effects of surface roughness initial stages of oxidation.24 The ultimate goal of this work was to improve life prediction for ZrB2-SiC of materials.  and  the  B. Fahrenholtz—contributing editor  Manuscript No. 33581. Received September 6, 2013; approved February 17, 2014.  Presented at  the 37th International Conference on Advanced Ceramics and Com posites, Ceramic Society, Daytona Beach, Florida, January 31, 2013.  *Member, The American Ceramic Society Member.  **Materials Advantage Student Member.  †  Author to whom correspondence should be addressed. e-mail: kns9a@virginia.edu  2279  J. Am. Ceram. Soc., 97 [7] 2279-2285 (2014)  DOI: 10.1111/jace.12911  © 2014 The American Ceramic Society  Journal  \\x0c', 'II.  Experimental Procedure  ZrB2-30 vol% SiC specimens were University of Science and Technology using attrition milled then hot pressed.25,26 WC contamina fabricated  at Missouri  powders which were (~2 wt%) was  tion  observed  due  to  the  attrition milling,  which used WC milling media  in polyethylene  jars. Speci mens were  cut  from a  series of billets using an automated 40 mm 9 4 mm 9 3 mm and  surface  grinder  into  bars  of  were finished using 1200 grit diamond abrasive. These bars (~7 mm 9 4 mm 9 3 mm)  were  cut  to size  for box furnace  testing using a diamond blade. Prior to heating, all specimens  were  cleaned  using  baths  of  detergent  and DI water, DI  water, acetone and then ethanol, each for 2 min in a sonica tor. Dimensions  and  initial weight  of  all  specimens were  recorded. Specimens were  individually oxidized and charac terized. The  specimens were placed in curved sections of an  yttria-stabilized zirconia (Ortech, Inc., Sacramento, CA)  tube  while in the furnace providing line contacts between the spec imen and boat, as  shown in Fig. 1(a). This boat helps pre vent  contamination  of  the  specimen  and  allows  for  easy  removal  from the  furnace. The weight of  the boat was also  measured before and after  testing and no significant weight  change was found. The tests were carried out in a box furnace (15 cm 9 15 cm 9 15 cm inner dimensions) with moly disilicide heating elements  (RapidTemp CM) under  stagnant  ambient air conditions. The furnace was brought  to tempera ture  and then the specimens were placed inside. When opened, the furnace temperature fell ~200°C, but returned to within 10°C of  the testing temperature in 10 s and was con stant  at  the  testing  temperature within 10 min. Tests were (1300°C, 1400°C, 1500°C,  run for 1550°C)  a  series of  temperatures  for  typically 100 min, using material  from multiple  bars. Timing of  each test was begun upon insertion of  the  specimen into the  furnace. Tests  conducted on material  cut  from the same bar are clearly indicated in the results. Speci mens were removed from the hot to room temperature (<5 min) were measured and recorded. Oxidized specimens were stored  furnace and cooled rapidly  in ambient air. Final weights  in  a  desiccator  cabinet  to minimize  reaction  of  any B2O3 remaining in the specimen with water vapor in the air. Other  tests were conducted at times as short 100 h at 1500°C. Two specimens were also oxidized in the box furnace for 100 min at 1500°C with the specimen placed at an ~45° angle, as shown in Fig. 1(b), in an attempt to  as  30 s  and up  to  characterize macroscopic glass flow.  The  oxidized  specimens were  characterized  by  Scanning  Electron Microscopy  (SEM-6700F;  JEOL,  Tokyo,  Japan)  and  Energy  Dispersive  Spectroscopy  (EDS;  Princeton  Gamma-Tech,  Inc.,  Princeton, NJ)  after  the  surface was  coated with a thin layer of carbon using the Precision Etch ing and Coating System. An operating voltage of 5 kV was  used to maximize the EDS signal  intensity for light elements.  Specimens  for  cross  section were mounted  in  epoxy while  under vacuum,  cross  sectioned using a diamond blade, and  remounted in epoxy. The first  epoxy mount was performed  in an attempt  to prevent pullout and spallation of  the oxide  during sectioning and polishing. The cross sections were polished using diamond polishing paste down to 1 lm in a mix ture  of  ethylene  glycol  and  200  proof  ethanol,  and  then  coated with  carbon. Nonaqueous  ethylene  glycol  and  200  proof ethanol were used as lubrication during specimen prep aration to minimize loss of B2O3, which is Average grain sizes were determined using the lineal  soluble in water.  intercept  method in at  least  three regions of each specimen and using  at  least  50  grains  per  area. Oxide  layer  thicknesses were  determined by measuring the thicknesses in a minimum of 20  locations over the cross section and averaging.  III.  Experimental Results  Oxidation Kinetics Variation  The specific weight change results after 100 min oxidation at each temperature are shown in Fig. 2. At 1500°C,  the weight  2.54 \\x06 0.73 mg/cm2 where gain for seven tests under dard deviation. When specimens 1.77 \\x06 0.26 mg/cm2. Average from a single bar were tests performed for 100 min are  identical  testing parameters was  the uncertainty reflects one  stan tested,  the weight  gain was  oxide  thicknesses over  all  reported in Table I. The  cross  sections  in Fig. 3 show that  even within a single  specimen large variation in oxide  layer  thicknesses  occurs.  Standard  deviations  of  the  oxide  layer  thicknesses are between 30% and 80% of  the average thick ness within any single specimen. The 1500°C gained 19.0 mg/cm2 24 h at and the dized for 100 h at 1500°C gained 45.3 mg/cm2.  specimen oxidized for  specimen oxi All oxidation weight gain results  (except  the 24 and 100 h  tests) were combined and plotted as weight gain per  surface  area versus  square  root of  time,  to determine  the parabolic  (a)  (b)  Fig. 1.  ZrB2-30 vol% SiC (YSZ) boat (a) for standard tests  specimen  in  yttria-stabilized  zirconia an ~45°  (scale  is  in cm)  (b)  at  angle.  Fig. 2.  ZrB2-30 vol% SiC weight gain after showing large variation in material  100 min exposures  in  air  behavior,  even  when  specimens are cut  from the same initial bar. Bar X indicates that no  specific bar was identified.  2280  Journal of the American Ceramic Society—Shugart et al.  Vol. 97, No. 7  \\x0c', 'oxidation rate  constant  and best account times.27 The  for  any  transient  oxidation  behavior  at  short  two  longer  term  tests were  excluded to analyze  the predictive  capabilities of  the results  from times of 4 h and less. From the plots, para bolic  rate  constants  (kp) were Table II. A 95% confidence interval was determined for  calculated  as  reported  in  the  slope of each plot, and from this, a maximum kp was calculated. The weight gain and calculated rate constants are 1500°C. As these specimens showed any indications of SiC-depletion beneath the oxide scale.28 Analysis of SiC-depletion behavior will be provided in a future paper.29  shown in Fig. 4  for  a point of note, none of  (A)  Viscous Flow of Glass:  The impact 1500°C for  of  oxidizing  the  specimen  at  an  ~45°  angle  at  100 min  is  shown in Fig. 5 as oxide  layer  thickness versus normalized  position. Measurements were taken for both the borosilicate  glass  thickness  and the ZrO2 the top surface of the specimen from the high end to the low  thickness  along  the  length of  end. As is seen in all other specimens,  there is a large scatter  in the thicknesses of  the oxide layers.  It  is  shown with 99%  certainty  that  no  correlation  between  oxide  thickness  and  specimen position exists,  indicating macroflow of  the borosil icate glass has not occurred under these exposure conditions.  (B)  Glass Pools  and Bubble Formation:  Observation  of  the  surface  on  specimens  after oxidation showed that large as 174 lm in diameter little as 30 s at 1500°C.24 An  pools of glass and bubbles as  formed on the  surface  in as  example of a pool and evidence of  the bubbles  for  the 30 s  exposure  condition are  shown in Fig. 6. These bubbles  are  dispersed  unevenly  over  the  specimen  surface  and  do  not  appear on all specimens. Table III summarizes observed bub ble size and distribution after various oxidation exposures. A  large  fraction of  specimens had no bubbles, demonstrating  difficulty in preserving bubbles in posttest specimens. Charac terization of specimen cross sections after oxidation in air for 10, 50, and 100 min at 1500°C showed 7 to 12 scallops of  differing widths and depths in the ZrO2 shown in Fig. 3. The specimen oxidized  layer, with examples  for  100 h  also  showed three scallops. As  the corners of  this  specimen were  Table I.  Thickness of Oxide Layers Formed in ZrB2 after Oxidation in Air for 100 min as a Function of Temperature  Temperature (°C)  ZrO2 thickness (lm)  Standard deviation (lm)  Borosilicate thickness (lm)  Standard deviation (lm)  Total oxide thickness (lm)  Standard deviation (lm)  No. of samples measured  (Avg. no. of measurements  per sample)  1300  25.1  8.7  7.7  6.0  32.8  6.3  1 (19)  1400  15.3  13.3  22.1  15.3  37.5  17.3  3 (35)  1500  21.9  15.6  27.6  20.5  39.6  26.4  6 (45)  1550  23.2  10.7  14.9  6.9  38.0  9.9  1 (64)  (a)  (b)  Fig. 3. Cross 1500°C  section of ZrB2-30 vol% SiC exposed for 100 min at and (b) 1400°C in air showing variation in layer  (a)  thicknesses and scalloped penetration depths of oxidation.  Table II.  Comparison of Literature and Experimental Results for Oxidation of ZrB2-SiC in Air  Paper  SiC content  Temperature (°C)  Time (min)  Weight change (mg/cm2)  kp (mg2/cm4 h)  Maximum kp  (95% confidence)  (mg2/cm4 h)  This study  30  1300  100  2.57 \\x06 0.4 3.36 \\x06 0.8 2.73 \\x06 0.8 3.14 \\x06 0.5 1.25  4.0  10.6  1400  100  7.4  24.7  1500  100  4.4  12.9  1550  100  4.4  11.6  18, 36  20  1327  10  4.9, 6.3  50  2  100  2.5  1627  10  3.75  100  50  8.75  100  13  37  20  1627  10  1.8  10.94  50  3  100  4.5  33.3  1627  10  3.5  NA  50  5.2  100  5.3  July 2014  Variability of ZrB2-30vol%SiC  2281  \\x0c', 'highly oxidized,  it  is possible more scallops had been present  earlier  in  the  test. The  bottom surfaces  of  all  specimens  showed fewer scallops.  IV.  Discussion of Results  Oxidation Kinetics Variability  It  is  important  to understand the sources of oxidation vari ability and the effect of each mechanism to accurately predict  lifetimes. It  is apparent  from the series of 100 min exposures  that within a batch of ZrB2-30 vol% SiC and even within one specimen bar, a large variation in oxidation behavior  exists. A loose  inverse  correlation between the  thickness of  the ZrO2 due to the  layer and the borosilicate  layer  is observed,  likely  impact of  the borosilicate  layer on oxidation, as  seen in Fig. 7. The borosilicate glass acts as an oxygen diffusion barrier, protecting the ZrB2-SiC, thicker less ZrO2 forms, whereas less borosilicate allows  so when the glass  is  for  Fig. 4.  Specific weight 1500°C,  change  (mg/cm2)  for  ZrB2-30 vol% SiC average kp and the 95% calculated from results up to 4 h.  oxidized  at  compared  to  both  confidence  interval on kp, corresponds to the boxed  Insert  region.  Star  indicates maximum weight  gain predicted using maximum oxidation depth.  Fig. 5. Oxide thickness versus position for a ZrB2-30 vol% SiC specimen oxidized at ~45° angle for 100 min at 1500°C in air.  (a)  (b)  (c)  Fig. 6.  ZrB2-30 vol% SiC after showing (a) a pool of borosilicate glass  30 s  exposure  at  1500°C in  air  (b) bubbles on the  surface,  highlighted with white,  and  (c)  the ZrO2  grains within  a  burst  bubble.  Table III.  Observations of Bubbles in Borosilicate Glass Phase Formed during Oxidation of ZrB2-30vol% SiC in Air  Temperature (°C)  Time  Bubble size (lm) Glass pool size (lm)  Additional  information  1500  30 s  43-174  —  Bubbles covered 34% of one region (54 cm2), 0% of another region  30 s  — —  13-50 61-79  No bubbles  1 min  No bubbles  100 min  977  — — — — — — —  1 bubble only  100 min  — —  No bubbles  100 min  No bubbles  100 h  10-395  2 bubbles evident  in cross-section  1550  10 min  — —  No bubbles  50 min  No bubbles  100 min  700  1 bubble only  2282  Journal of the American Ceramic Society—Shugart et al.  Vol. 97, No. 7  \\x0c', 'greater growth of ZrO2. This has also been observed by Carney et al.,19 and will be discussed in more detail when bubble  formation is addressed. While  some variability in oxidation  kinetics between different specimens may be due to specimen  preparation and technique or  variation in oxidation condi tions  (relative humidity,  impurities),  the mechanism for  sig nificant amounts of variability within an individual specimen  remains unexplained.  The specific weight gain which would occur if the maximum  oxidation depth (star in Fig. 7) were achieved over the entire specimen when oxidized at 1500°C for 100 min was calculated.  Several assumptions were made. The ZrO2 which forms was assumed to be a fully dense layer. The consumption of ZrB2 and SiC was assumed equal to the starting mole fraction. The  SiO2 was Twenty-three percent of  assumed  to  form a  layer  on  top  of  the ZrO2. the B2O3 formed was assumed to be retained in the glass, whereas the remainder was assumed to  vaporize. This value was based on results in which the compo sition of the glass formed on a ZrB2-30 vol% SiC specimen during oxidation in air at 1500°C for 100 min was determined  using Inductively Coupled Plasma Optical Emission Spectrometry.30 The weight gain due to oxygen incorporation in  condensed oxide phases and the B and C mass  loss due  to  B2O3 and CO vaporization was calculated per Zr atom oxidized. The number of moles of Zr in the oxidized volume \\x004. (96.5 lm 9 1 cm 9 1 cm) was calculated to be 4.427 9 10 This leads to an estimated maximum specific weight gain of 14.3 mg/cm2 for 100 min oxidation at 1500°C, star in Fig. 4. This predicted weight gain is almost as high as the weight gain after 24 h at 1500°C and will be discussed  shown by the  below.  Included in Table II are literature results for weight gain during oxidation for ZrB2-SiC, and measured kp. Most studies do not include sufficient quantitative results from the oxi dation  tests, making  comparison  between  studies  difficult;  however,  as  seen  in Table II  the  rate  constants measured  here are of  the same order of magnitude as those in the liter ature.  The 24 and 100 h oxidation weight gains were compared to predictions from tests conducted for times ≤4 h in Fig. 4. The experimental weight gains of the 24 and 100 h specimens were  both greater  than the upper 95% confidence interval on kp, indicating that use of short term oxidation tests (the literature  focuses on 10 min to 4 h) is insufficient for prediction of long  term results.  It  is possible  that  the difference between long  duration weight gain and that predicted by shorter term tests  indicates a change in mechanism. However,  the morphology  of  the  microstructure  (grain  shape,  oxide  distribution,  presence of  scallops) appears  to remain the same after  long and short-term tests. The 24 and 100 h specimen weight gains  both fall above the rate predicted by the maximum oxidation  depth and the upper 95% confidence interval calculated from  weight  change. These  results  reinforce  the need to consider  maximum oxidation ZrB2-SiC.  depth  for  life  prediction modeling  of  (A)  Viscous Flow of Glass:  From the specimen oxidized at an ~45° angle, measureable macroscopic flow of the borosilicate surface at 1500°C. Surface  results  of  the  it  can be  seen that no  glass  occurred along the  tension must  prevail over gravity in determining the flow of  the borosili cate  glass  and  viscous  flow  on  a macro  scale is not 1500°C.  responsible  for  variable  oxidation  kinetics  at  It  should  be  noted,  however,  that  viscosity  of  the  glass will  decrease with increasing  temperature  and that  viscous flow  may influence above 1500°C.19  the  oxidation  behavior  at  temperatures  (B)  Bubble Formation in Glass Scale:  Bubble  forma tion due to generation of gaseous oxidation products  is  sug gested  as  the most  likely  source  of  oxidation  variability.  Formation of bubbles  in the  glassy phase  could be due  to  CO(g)  formed  during  SiC  oxidation  [Reaction (2)]  or  B2O3(l) tion (3)]. FactSage, a computer program for chemical thermodynamics calculations,31 was used to calculate the partial pressures of the gasses which form upon oxidation of ZrB2- 1500°C. The equilibrium constant for Reaction (1), 2.022 9 1039.  volatilizing  due  to  its  high  vapor  pressure  [Reac SiC at  determined  using  the  Reaction module  is  Using:  Keq ¼ aZrO2 aB2O3 ðlÞ aZrB2 P O2  5 2  (4)  and  assuming  that  the  activity of  all  solids  and  liquids  is  unity,  the equilibrium partial pressure of oxygen at the ZrB2/ \\x0016 atm. Using calculated to be 1.9 9 10 inputting ZrB2-SiC, and holding the equilibrium value of PO2, partial pressure of other gasses formed were calculated. The  ZrO2 the Equilibrium module,  interface was  the activity of oxygen at  the  three dominant \\x001 atm, gas partial pressures were \\x005 atm, as PCO = 9.97 9 10 PCO2 = 9.07 9 10 \\x003 atm. At PSiO = 3.25 9 10 1500°C, B2O3(g) \\x003 atm, has a pressure of 2.6 9 10 independent of PO2. CO(g) has the highest partial pressure and with a partial pressure of  follows:  and  partial  almost 1 atm,  it  is  reasonable  that CO(g) will  form bubbles  in the borosilicate glass.  Figure 8 illustrates a mechanism by which oxidation behav ior is affected by bubble formation. Results from initial oxida tion behavior described in a companion paper are also used to inform the hypothesis.24 In this mechanism, borosilicate glass  covers the surface after 30 s, at which point bubbles begin to  form. When  bubbles  burst,  exposing  underlying material,  increased oxidation is allowed in the region. The mechanism  presented in Fig. 8 is consistent with observations of Gangireddy et al.32 for ZrB2-15 vol% SiC material oxidized between 1450°C and 1650°C and Carney et al. for ZrB2-20 vol% SiC material oxidized for 150 min at 1600°C.19 The formation and  bursting of bubbles was attributed there to the pressure of the  gaseous products and viscosity of the glass layer. Carney et al.  report  that no bubbles were  formed on the bottom surface,  which they say contributes  to the uniformity of  the bottom  oxide scale. The data reported in Fig. 7 support the proposed  mechanism. When the borosilicate  layer  is  thick, underlying  layers of ZrO2 are provided by the silica-rich glass. When the borosilicate layer is  thin, due  to greater oxidation resistance  thin, the region shows a large range of ZrO2 layer thicknesses. This large distribution of ZrO2 thicknesses with thin glass layers suggest bubble generation and rupture over a range of oxi dation times.  Fig. 7. Dependence of ZrO2 thickness on borosilicate thickness for a specimen oxidized 100 min at 1500°C in ambient air. Star indicates  maximum oxidation depth, used in maximum weight gain prediction.  July 2014  Variability of ZrB2-30vol%SiC  2283  \\x0c', '(C) Summary: Viscous  glass flow and bubble  formation  have been considered as possible  explanations  for observed  variability in oxidation kinetics  for ZrB2-30 vol% SiC. Viscous flow has clearly been ruled out as a major contributor  to this variability. Surface roughness effects are described in  a companion paper and shown to be unimportant 30 s.24  for  times  greater  than  Bubble  formation  is  the most  likely  mechanism for variability in oxidation behavior as illustrated  in Fig. 8.  Limitations  to  this  study  and  the  proposed mechanism  include the following. This  study was conducted on bars cut  from a series of hot-pressed billets; variations due to process ing differences were not  identified. However,  the variation in  results observed herein are similar literature.13,33-35 Viscous flow of be more important as gas velocities  to those  reported in the  the glass phase  is  likely to  and temperatures  rise.  Different volume  fractions of SiC will  result  in more  silica rich  scales,  potentially  affecting  oxidation  results. Bubbles  were not observed in the glass phase of all  specimens  (Table  III), however, previous work has  shown bubbles  form and  burst in as little as 1 min, making them difficult to preserve in posttest specimens.32 Further study of bubble formation is  required  to  confirm this  proposed mechanism for  a wider  range of materials and oxidation conditions.  Prior  to use of  these materials,  an oxidation life predic tion model  is needed. Previous experimental work forces oxi dation kinetics to a parabolic fit as oxidation is assumed to a diffusion limited process.36,37 Also, most  be  experimental  work also focuses on oxidation times of 10 min to 2 h.  In  addition,  current models  of  oxidation  behavior  predict  an  average depth of oxidation, or glass time.21,23 This work shows  thickness, as a function  of  that using the average value  for weight gains obtained for  times of 4 h and less  to deter mine a 95% confidence  interval on the parabolic oxidation  rate constant  is not sufficient  for predicting oxidation behav ior at  longer  times. First,  there are regions where the depth  of attack is much greater  than average and,  second,  short term oxidation tests  are not  good predictors  for  long-term  tests. Consideration of a maximum attack rate  is needed to  conservatively  determine  the  capability  of  this material  to  survive  exposures  to  ultrahigh  temperature  oxidizing  envi ronments.  V.  Conclusions  ZrB2-30 vol% SiC showed variable oxidation kinetics. The temperatures between 1300°C specific weight gain in air at and 1550°C and exposure  times between 30 s  and 100 min  varied up to a factor of 2.4, even when well-controlled test ing  parameters  are  used. The  oxide  thickness  formed  on  any  one  specimen  showed  variations  in  thickness  up  to  80% of the average (≤4 h) under-predict emphasizing the need  thickness.  Short  duration  experiments  the weight  gain at  long  times  (100 h),  for  better  life  prediction methods.  Variable  oxide  thickness  and  oxide  phase  distribution  formed on the  surface of  the  specimen likely explains  scat ter  in measured weight  change  among different  specimens.  Viscous glass flow has been eliminated as a primary cause this variability at 1500°C. A mechanism based on bub for  ble  formation in the glass phase  is proposed to explain the  observed oxidation variability. During  initial  stages of oxi dation, borosilicate glass  spreads unevenly over  the  surface  of  the material,  slowing oxygen diffusion in some areas  (i.e.,  where  it  is  thicker) more  than other areas  (i.e., where  it  is  thinner or only ZrO2 glass occurs predominantly  is present.) Bubble  formation in the  due  to  accumulation  of CO(g)  produced  by  oxidation.  The  bubbles  burst  leaving  areas  with  thin  or  no  borosilicate  coverage;  thereby  further  increasing  inhomogeneity  in  the  oxide  scale with  resulting  variability in oxidation kinetics.  Acknowledgments  The authors would like to acknowledge Eric Neuman and Dr. William Fahr enholtz at Missouri University of Science and Technology for the ZrB2-30 vol % SiC material. Initial funding for this work was provided by The National  Hypersonic Science Center-Materials and Structures.  References  1A. Paul, D. D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. G. P. Bin ner, B. Vaidhyanathan, A. Heaton, P. 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  "_id": 127,
  "PDF": "Microstructural characterization of ZrC-MoSi2 composites oxidized in air at high temperatures.pdf",
  "Text": "['Applied Surface Science 283 (2013) 751- 758  Contents lists available at SciVerse ScienceDirect  Applied Surface Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / a p s u s c  Microstructural characterization of ZrC-MoSi2 composites oxidized in air at high temperatures  Ludovic Charpentier a,∗ , Marianne Balat-Pichelin a , Eric Bêche a , Diletta Sciti b , Laura Silvestroni b  a PROMES-CNRS,  7  rue  du  four  solaire,  66120  Font-Romeu-Odeillo,  France  b ISTEC-CNR,  Via  Granarolo,  48018  Faenza,  Italy  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article  history:  Received 21 March 2013 Received in revised form 2 July 2013 Accepted 4 July 2013 Available online 12 July 2013  Keywords:  Ceramic SEM XPS High temperature corrosion Internal oxidation  Future receivers of concentration solar power plants have to heat the coolant (air) so that its temperature will exceed 1300 K. Such absorbers require materials able to support thermal and mechanical stresses, with the slowest possible oxidation kinetics at very high temperatures. Zirconium carbide (ZrC) with silicon additive like molybdenum disilicide (MoSi2 ) could be a candidate material for such an application. ZrC/20 vol% MoSi2 samples were oxidized for 20 min at the 5 kW Odeillo solar furnace in air at various temperatures between 1800 and 2000 K and the oxidation behaviour was investigated as a function of the exposure temperature and the surface machining preparation. X-Ray photoelectron spectroscopy and scanning electron microscopy coupled with energy dispersive X-Ray spectroscopy enabled to study the microstructure evolution and to identify the oxidation mechanism leading to the formation of zirconia and silica layers. Based upon the characterizations, we can affirm that ZrC/20 vol% MoSi2 seems able to withstand temperatures up to 2000 K in air.  © 2013 Elsevier B.V. All rights reserved.  1.   Introduction  Inside  the  current  concentration  solar  power  plants,  the receivers  in metallic alloys do not enable  the completion of  the heating of  the coolant air up  to 1300 K due  to  the severe oxidation of these materials beyond 1200 K [1]. Ultra-high temperature ceramics  (UHTC) possess good  thermal stability and mechanical properties up to 1800 K and recent studies have also shown that they have more favorable radiative properties [2-5], as compared to silicon carbide. Thus, they can be considered as an alternative to conventional metallic materials enabling a notable  increase of the working  temperature and  thus of  the conversion efficiency [2-5]. Among potential UHTC, zirconium carbide  (ZrC) possesses a high melting point (3500 K) and interesting mechanical properties, especially a high hardness of 27 GPa  [6]. The main  limitation  for its high temperature applications is the oxidation kinetics. Porous zirconia  (ZrO2 )  is non-protective and ZrC continuously oxidizes according to a linear kinetics even at temperatures as low as 800 K [7]. Incorporation of silicon was found to enable the formation of mixed oxide layer such as zircon (ZrO2 ·SiO2 ), which was shown to  ∗ Corresponding author. Tel.: +33 468 307 741; fax: +33 468 307 799x302 940.  ludovic.charpentier@promes.cnrs.fr (L. Charpentier).  address:  E-mail  0169-4332/$ - see front matter ©   2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2013.07.012  Pression  et  Hautes  Température  Solaire) facil improve the resistance to oxidation [8-11]. This incorporation can be achieved by sintering ZrC with molybdenum disilicide (MoSi2 ) additive, using spark plasma sintering (SPS) [12] or even by pressureless sintering upon addition of a suitable amount of MoSi2 [13,14]. In a preliminary study, a pressureless sintered ZrC/20 vol% MoSi2 material was oxidized  in  the  temperature range 1800-2400 K  in the REHPTS (Réacteur ity at the 5 kW Odeillo solar furnace  in flowing air [15] and some surface analyses revealed severe surface degradations with bubbles bursting above 2000 K. Nevertheless, this study needed deeper investigations  to understand how  the oxide  layers evolved with temperature and how the initial surface state may affect their compositions. Therefore complementary analyses have been performed on samples oxidized at 1800 K and 2000 K. X-Ray photoelectron spectroscopy  (XPS) enabled  to analyze  the  chemical bonds present close  to  the  surface and  to evidence  the presence of  silica and zirconia bonds, and some distinctiveness due  to  the  initial surface  state of  the  sample and  to  the  temperature. Cross-section scanning electron microscopy (SEM) coupled with energy dispersive X-Ray spectroscopy (EDS) has given clearer indication on the oxide  layer composition. These more complete characterizations enabled us to suggest an oxidation mechanism for this composite material.                                                                            \\x0c', '752  L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  Fig. 1. SEM images of reference samples machined showing (a) the polished surface, by the as-machined surface by (b) EDM or (c) DLT.  2. Material and methods  2.1.  Material  The  material  investigated  was  a  pressureless  sintered ZrC/20 vol% MoSi2 produced with  the  following  commercial powders:  ZrC  (initial mean  grain  size:  3.8  \\u242em,  supplier: H.C. Starck)  and MoSi2 (initial mean  grain  size:  2.8  \\u242em,  supplier: Aldrich). The mixture was ball milled  in ethanol, dried by rotary evaporator, sieved  through 250  \\u242em screen and green shaped  to form a 35 mm pellet by applying first  linear pressure and  then cold  isostatic  pressure. Afterwards,  the  cylindrical  pellet was sintered in a graphite furnace at 2220 K for 60 min in flowing argon without applied  load. A relative density of 95% was measured by Archimedes’ method. SEM, TEM and XRD analysis have shown the presence of faceted ZrC  grains with dimensions up  to 6  \\u242em, homogeneously  surrounded by MoSi2 phase, which displays an  irregular shape with low dihedral  angles,  as  Fig. 1(a)  shows. Minor phases  are SiC grains, around 1 vol% and Zr-Si phases with various stoichiometries [13,16]. The samples (Ø  = 25 mm,   = 2 mm) were obtained by electrical discharge machining (EDM) from the sintered cylindrical pellet that had an average density of 6.18 g cm−3 . EDM technique consists  in material removal by a series of rapid current discharge between two electrodes, the material locally melts and, on the surface, local formation of zirconia can occur. Some samples were further machined using a Diamond Loaded Tool (DLT), which enabled the removal of the oxide  layer  formed during EDM. From now on,  the polished samples will be referred to as DLT samples, the non-polished samples as EDM samples. The surface roughness Ra  (Taylor Hobson Talysurf Plus profiler) was around 0.13  for DLT samples and 0.6  \\u242em for EDM samples. Differences  in surface aspect are  reported  in previous publications  [15,16], but  for  the sake of clarity  the main  features are summarized. Samples machined by EDM (Fig. 1(b)) have a darker appearance  compared  to  the DLT  and  SEM  images  revealed  a rougher surface for EDM samples and areas where ZrO2 was present (Fig. 1(b)), whilst DLT samples showed some grain pull-out due to the diamond polishing (Fig. 1(c)).  \\u242em   h  2.2.  Experimental  setup  ±  achieved at very high heating rate (up to 100 K s−1 ) on a homogeneous 10 mm diameter area. Two mirrors enable a monochromatic (5  \\u242em) optical pyrometer to measure the surface temperature of the sample through a fluorine window. The accuracy of the temperature measurements goes from 1400   15 K to 2100   22 K. As the formation of a zirconia layer occurs as soon as the experiments start,  the normal spectral emissivity of  the material was  that of zirconia, reported to be 0.75 at 5  \\u242em [15]. The oxidations were performed in flowing air. One entrance of the reactor was open  in order  to enable a constant renewing of the atmosphere surrounding the sample. Due to the altitude of the laboratory, the total atmospheric pressure is around 87 kPa and the oxygen partial pressure pO2 is 17 kPa. The surface temperature of the samples was maintained at a constant plateau during 20 min oxidation.  ±  2.3.  Post-experimental  characterisations  (h(cid:2)  XPS analyses were performed using a SIA Riber Cameca UHV 10−6 Pa.  device  operating  at  a  pressure  of  The  photoelectron emission spectra were  recorded using an Al-K␣  = 1486.6 eV) non-monochromatic source (600 W). The analyzed area was about 25 mm2 . The surface contaminant (atmospheric carbon) must be removed by ion sputtering of the surfaces (Ar+ ion beam accelerated at 3 keV). The kinetic energy of the photoelectrons was measured using a Riber Cameca MAC 2 spectroscopic two stages spectrometer. The analyser resolution was fixed at 1 eV. Spectrometer energy calibration was made using the Au 4f7/2 (83.9   0.1 eV) and Cu 2p3/2 (932.8   0.1 eV) photoelectron  lines. XPS spectra are recorded  in direct N(Ec). The background signal was removed using the Shirley method [20]. The crystalline phases  formed after oxidation were detected by X-ray diffraction  (Siemens D500, Karlsruhe, Germany). The microstructure evolution was analyzed using  scanning electron microscopy (SEM, Cambridge S360, Cambridge, UK) coupled with energy dispersive  spectroscopy  (EDS,  INCA Energy 300, Oxford instruments, UK) on the surface and cross-section to estimate the thicknesses and compositions of the various oxide layers.  ±  ±  3. Results and discussion  3.1.  XPS  surface  analyses  The behaviour of UHTC  in oxidizing atmosphere at very high temperatures was studied using the REHPTS facility [17-19]. A flat mirror (heliostat), whose position  is servo-controlled to the apparent movement of the sun, reflects the incident solar flux to a concentrator with facetted mirrors. A shutter enables to control the  fraction of the concentrated solar flux delivered to the sample placed  inside the reactor and therefore to monitor  its surface temperature. In this set-up, the sample is placed 25 mm above the focus of the solar  furnace, so that very high temperatures can be  The samples studied have produced charging shifts. The binding energy scale was established by referencing the C 1s value of adventitious carbon (284.6 eV) [21]. The C 1s and O 1s photoelectron peaks were analyzed by Gaussian/Lorentzian peak fitting. The full width at half maximum (FWHM) and the positions of the components are similar to those collected on our reference samples. The Zr 3d3/2,5/2 , Si 2p1/2,3/2 and O 1s spectra collected  for EDM and DLT  samples are  shown  in Figs. 2 and 3,  respectively. The spectra were deconvoluted with 4 or 2 components.  Zr 3d3/2,5/2                                  \\x0c', 'L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  753  Fig. 2. Zr 3d3/2,5/2 , Si 2p1/2,3/2 and O 1s spectra collected for EDM Zr/20 vol% MoSi2 samples before and after oxidation during 20 min at 1800 and 2000 K in air.  ±  The ratio of the peak  intensities Zr 3d5/2 /Zr 3d3/2 was fixed close to  the  theoretical value of 1.50   0.05. The assignment of  the Zr 3d3/2,5/2 , Si 2p1/2,3/2 and O 1s components  from  the XPS spectra collected  for  the different  samples are  summarized  in Table 1. The Zr 3d3/2,5/2 peak positions  located at 181.2 eV (Zr 3d3/2 ) and 179.0 eV  (Zr 3d5/2 ) are characteristic of Zr-C bonds  in ZrC compound [22-24]. The Zr 3d3/2,5/2 peak positions located at 184.9 eV (Zr 3d3/2 ) and 182.6 eV (Zr 3d5/2 ) are attributed to Zr-O bonds  in ZrO2 compounds (Zr-O4 environments, Zr4+ states) [21-23,25,26].  The Si 2p1/2,3/2  spectra were deconvoluted with 2 or 1 components. The  components  located  at 103.2  and 99.6   0.1 eV  are respectively characteristic of Si-O bonds  in SiO2 compound and Si-Mo bonds  in MoSi2 compounds [27,28]. The O 1s spectra were deconvoluted with 2 components, with an observed shift about 1.6 eV. The component located at 530.7   0.1 eV is attributed to OZr bonds  in ZrO2 compounds  [22,24,25]. The component  located at 532.3   0.1 eV  is attributed  to O-Si bonds  in SiO2 compounds [17].  ±  ±  ±  Table 1 Assignment of the Zr 3d3/2,5/2 , Si 2p1/2,3/2 and O 1s components from the XPS spectra collected.  Spin-orbit doublet   ±   0.1 eV   B.E.   FWHM   ±   0.05 eV   Chemical bonds   Compounds  Zr 3d5/2  Zr 3d3/2  Si 2p1/2,3/2  O 1s  179.0  182.6  181.2  184.9  99.6 103.2  530.7  532.3   2.0  2.3  2.0  2.3  2.2  2.4  2.3  2.3   Zr-C  Zr-O  Zr-C  Zr-O  Si-Mo  Si-O  O-Zr  O-Si   ZrC ZrO2 ZrC ZrO2 MoSi2 SiO2 ZrO2 SiO2  Bold and underlined element is the atom the analyzed electron is belonging to, the other is the atom the analyzed electron is linked to.                \\x0c', '754  L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  Fig. 3. Zr 3d3/2,5/2 , Si 2p1/2,3/2 and O 1s spectra collected for DLT Zr/20 vol% MoSi2 samples before and after oxidation during 20 min at 1800 K in air.  In the reference unoxidized samples (Fig. 2(a) and Fig. 3(a)) Zr-C bonds and Si-Mo bonds are evident, but  in all the oxidized specimens Zr-C bonds are no longer detected. The native oxide that has grown during EDM treatment (reference  in Fig. 2(a)) seems to be a mixture of zirconia and silica phases (presence of Zr-O and Si-O bonds). ZrO2 is noticeably reduced after DLT polishing but a new SiO-based one has grown  (reference  in Fig. 3(a)). Comparing ZrO/Zr-C and Si-O/Si-Mo ratios corresponding to Fig. 2(a) and Fig. 3(a), it seems that the native oxide on DLT sample is richer in silica than the native oxide on EDM sample. The DLT sample oxidized at 1800 K (Fig. 3(b)) contains a lower amount of Si-Mo bonds as compared to the reference,  indicating that some Mo-Si phase is still present after oxidation. On the other hand, in the EDM one, no Mo-Si phase can be detected (Fig. 2(b)). The O 1s spectra display different proportions of oxygen bonds, which  indicates different composition of the oxide as a  function of the surface finishing,  in Fig. 2(b) and Fig. 3(b), and of the temperature,  in Fig. 2(b) and  (c). As  far as  the machining method  is concerned, it appears that the DLT polishing favours the formation of Zr-rich phases over Si-rich ones, whereas EDM sample presents equivalent amount of Zr-rich and Si-rich phases after oxidation at 1800 K  for 20 min. Whilst as  for the temperature, oxidation at 2000 K induces the formation of higher amount of SiO-based scale over ZrO2 (Fig. 2(c).  3.2.  Microstructure  evolution  The appearance of the discs after oxidation is depicted in Fig. 4. It can be observed that both samples oxidized at 1800 K have similar greyish aspect with smooth surface, whilst increasing the temperature to 2000 K induced a colour change to white-yellow, especially in the centre, where bubbling and little spalling occurred. X-ray diffraction revealed the formation of monoclinic ZrO2 and traces of MoSi2 and Mo5 Si3 on  the samples oxidized at 1800 K, whilst on the sample oxidized at 2000 K, only zirconia was present (see Table 2). Figs. 5-7 show cross-section SEM images of ZrC/MoSi2 samples that were oxidized for 20 min in air at 1800 and 2000 K. A summary  in  \\u242em   \\u242em   \\u242em   is reported   of the microstructural details for the three samples  Table 2. The microstructural modifications on the EDM sample at 1800 K can be divided into three layers (Fig. 5(a)), starting from the surface. The outermost oxide layer (Fig. 5(b)) is 170  \\u242em thick and it is composed of porous ZrO2 immersed in SiO2 , which is in agreement with the corresponding XPS spectra of Fig. 1(b) evidencing both Zr-O and Si-O bonds close to the surface. Underneath this scale, a 90  \\u242em thick layer based on porous Zr-O-C is present (Fig. 5(c)), thin SiO2 \\u242em film and large cavities are also present. Fig. 5(d) shows the 140  thick third layer where ZrC grains are surrounded by a Zr-O-C phase having darker contrast (see the EDS spectra in the inset) with some small pockets of oxidized MoSi2 . The oxide  layer of DLT  sample at 1800 K presents a different morphology (Fig. 6(a)). The outermost oxide scale (Fig. 6(b)) is 180  thick  and  contains mainly  Zr-O-C  and MoSi2 with its oxidation products,  i.e. SiO2 and Mo5 Si3 , supporting  the XPS analyses  that have  revealed  the presence of Zr-O, Si-O and SiMo  bonds  close  to  the  surface.  The  underneath  oxide  layer, magnified  in  Fig. 6(c)  and  (d),  is 200  thick  and  is mainly composed by ZrC grains enclosed by a Zr-O-C phase. At  triple junctions, MoSi2 surrounded by SiO2 and Mo5 Si3 can be observed (Fig. 6(d)). The structure of the oxide layer grown on EDM sample at 2000 K \\u242em thick SiO2 layer (Fig. 7(a))  is even more complex. A wavy 30  lies on the top of the specimen (Fig. 7(b)),  in agreement with the higher O-Si peak of the XPS profile of Fig. 2(c). Underneath the silica layer, characterized by dark contrast, about 150  thick brainlike ZrO2 is immersed in SiO2 , Fig. 6(b). Moving further toward the bulk, one meets 85  \\u242em thick partially porous Zr-O-C layer filled by SiO2 (Fig. 7(c)). The oxide  layer underneath  (Fig. 7(d))  is around 80  \\u242em thick and  it  is composed by dense ZrC grains surrounded by granulous Zr-O-C.  In  this region, both MoSi2 and Mo5 Si3 are present and are  interconnected by a SiO2 thin film, recognizable by the dark contrast. Right above the bulk, another layer is 90  \\u242em thick (Fig. 7(e) and (f)) composed with the same chemistry as above but with SiC original grains and higher amount of MoSi2 instead of Mo5 Si3 .                    \\x0c', 'L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  755  Fig. 4. Appearance of the discs after oxidation, compared to reference EDM and DLT samples. (For interpretation of the references to color in the text, the reader is referred to the web version of the article.)  3.3.  Oxidation  mechanisms  Based upon the cross-section analyses and XPS characterization, the following oxidation mechanisms can be proposed. The presence of Zr-O and Zr-O-C phases close  to  the  interface with the carbide show that the oxidation mechanism should  be similar to the one proposed by Shimada [29]. During the first stage of  the oxidation,  the Zr-O-C phase  is  formed  according to the reaction:  ZrC(s) +  xO2(g) →  ZrC1-xOx(s) +  xCO(g)  (1)  Table 2 Summary of the phases detected by XPS, XRD and SEM-EDS on the reference and oxidized samples.  Finishing   As sint. EDM  As sint. DLT  EDM  DLT  EDM  T   (K)   -  -  1800  1800  2000  XPS   ZrC, ZrO2 , MoSi2 , SiO2  ZrC, ZrO2 , SiO2 , MoSi2  ZrO2 , SiO2  ZrO2 , SiO2 , MoSi2  ZrO2 , SiO2  XRD   c ZrC t MoSi2 c ZrC t MoSi2 m ZrO2 t MoSi2  m ZrO2 ,  t MoSi2  m ZrO2  Thickness (\\u242em) and composition  -  -   170  90  140  180  200  30  150  85  80  90  -  -  SiO2 -ZrO2 SiO2 -ZrOC ZrC-ZrOC-MoSi2  Total  thickness:  400  ZrOC-SiO2 -Mo5 Si3 -MoSi2 ZrC-ZrOC-SiO2 -Mo5 Si3 -MoSi2  Total  thickness:  380  SiO2 SiO2 -ZrO2 SiO2 -ZrOC ZrC-ZrOC-SiO2 -Mo5 Si3 -MoSi2 ZrC-ZrOC-SiC-SiO2 -Mo5 Si3 -MoSi2  Total   thickness:  435                                  \\x0c', '756  L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  Fig. 5. Cross-section SEM images of an EDM ZrC/20 vol% MoSi2 samples after 20 min oxidation in air at 1800 K. (a) Global cross-section, (b) zoom in oxide layer 1, (c) zoom in oxide layer 2, (d) zoom in oxide layer 3 with EDS spectra of ZrC and Zr-O-C phases.  Then this phase  phase: ZrC1-xOx(s) +  (3/2-x)O2(g) →   ZrO2(s) +   (1-x)CO(g)  The main difference with Shimada model  [29]  is  that we did not observe any presence of free carbon, but Shimada has worked  is saturated with oxygen and evolves to ZrO2  (2)  in lower oxygen partial pressure (10 kPa) and temperature (870 K) than we did. The porous nature of ZrO2 probably enables gaseous CO to escape without significant cracking or spalling of the material. The difference between DLT  and EDM  samples oxidized  at 1800 K suggests that a smoother surface (that of DLT) slowers the oxidation kinetics and  thus  formation of  the  intermediate oxide  Fig. 6. Cross-section SEM images of a DLT ZrC/20 vol% MoSi2 samples after 20 min oxidation in air at 1800 K. (a) Global cross-section, (b) zoom in oxide layer 1, (c) zoom in oxide layer 2, (d) oxidizing MoSi2 in layer 2.                \\x0c', 'L.  Charpentier  et  al.  /  Applied  Surface  Science  283 (2013) 751-  758  757  Fig. 7. Cross-section SEM images of an EDM ZrC/20 vol% MoSi2 samples after 20 min oxidation in air at 2000 K. (a) Global cross-section, (b) zoom in oxide layer 2, (c) zoom in oxide layer 3, (d) zoom in oxide layer 4, (e) zoom in oxide layer 5, (f) magnification in layer 5.  through reaction (1) is more evident. This hypothesis is deduced by the higher amount of Zr-O-C phases inside the oxide layers grown on DLT samples (Fig. 5) and by a more important proportion of O-Zr bonds on the surface (Fig. 2(b)). As  for  the oxidation of  the secondary phase MoSi2 , SiO2 and Mo5 Si3 phases may  form due to the  following reaction, as previously studied by Wirkus [30]: MoSi2(s) +   7O2(g) →  Mo5 Si3(s) +  7SiO2(s)  (3)  SiO2 may also be produced  together with  gaseous molybdenum oxide, according to: MoSi2(s) +  2MoO3(g) +   7O2(g) →   4SiO2(s)  the   formation of  (4)  The newly formed SiO2 diffuses to the surface of the sample and this diffusion  is  further on  favoured at 2000 K by the  lower glass viscosity.  proved that the formation of zirconia passes through the formation of an intermediate Zr-O-C phase. The evolution of  the oxide  structure was  studied as a  function of  the  temperature, 1800 K and 2000 K, and of  the surface machining procedure, EDM and DLT. It seems that a smoother and less oxidized surface, like the one obtained by DLT, favours the formation of the Zr-O-C phase at 1800 K. The oxidation of MoSi2 at 2000 K enabled the formation of low viscosity silica that migrated to the surface. This glassy scale and the complex architecture of the oxidized  layer showed to be effective  in the protection of further oxygen inward, as the total thickness of the modified layer was just 35  \\u242em higher as compared to the sample oxidized at 1800 K. We can thus affirm that ZrC/20 vol% MoSi2 seems able to withstand temperatures up to 2000 K  in air. Complementary tests with cycling and  longer temperature exposures (ageing) will be performed to check if this material could effectively support the conditions of a solar receiver, like intermittent solar radiation and clouds.  4. Conclusions  References  The oxidation behaviour of a ZrC/20 vol% MoSi2 composite was studied in the REHPTS facility above 1800 K, in order to understand if  this ceramic could have an application  in concentrating solar power system plants, as absorber material. SEM and XPS analyses  [1] P.  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Shirley, High-resolution X-ray photoemission spectrum of  bands of gold, Phys. Rev. B 5 (1972) 4709-4714. [21] D. Briggs, M.P. Seah, Practical Surface Analysis, Auger and X-Ray Photoelectron Spectroscopy, John Wiley & Sons, New York, 1993. [22] M. Balaceanu, M. Braic, V. Braic, A. Vladescu, C.C. Negrila, Surface chemistry of plasma deposited ZrC hard coatings,  J. Optoelectron. Adv. Mater. 7 (2005) 2557-2560. [23] D. Craciun, G. Socol, N. Stefan, G. Bourne, V. Craciun, Chemical composition of ZrC thin films grown by pulsed laser deposition, Appl. Surf. Sci. 255 (2009) 5260-5263. [24] S. Kato, K.I. Ozawa, K. Edamoto, S. Otani, Photoelectron spectroscopy study of the oxidation of ZrC(1 0 0), Jpn. J. Appl. Phys. 39 (2000) 5217-5222. [25] D.Y.  Cho,  H.S.  Jung,  C.S.  Hwang,  Structural  properties  tronic  structure  of  HfO2 -ZrO2 composite  films,  Phys.  Rev.  http://dx.doi.org/10.1103/PhysRevB.82.094104. [26] C. Monticelli, F. Zucchi, A. Pagnoni, M. Dal Colle, Corrosion of a zirconium diboride/silicon carbide composite in aqueous solutions, Electrochim. Acta 50 (2005) 3461-3469. J. Meng, J. Lu, J. Wang, S. Yang, Preparation and properties of MoSi2 composites reinforced by TiC, TiCN, and TiB2 , Mater. Sci. Eng. A 396 (2005) 277-284. [28] M. Nayak, G.S. Lodha, R.V. Nandedkar, S.M. Chaudhari, P. Bhatt,  Interlayer composition  in Mo-Si multilayers using X-ray photoelectron spectroscopy,  J. Electron Spectrosc. Relat. Phenom. 152 (2006) 115-120. [29] S. Shimada,  Interfacial  reaction on oxidation of carbides with  carbon, Solid State Ionics 141-142 (2001) 99-104. [30] C.D. Wirkus, D.R. Wilder, High temperature oxidation of molybdenum disilicide, J. Am. Ceram. Soc. 49 (1966) 173-177.  elec(2010),  and  B   formation of                \\x0c']"
},{
  "_id": 128,
  "PDF": "Microstructure evolution of a W-doped ZrB2 ceramic upon high-temperature oxidation.pdf",
  "Text": "['O R I G I N A L A R T I C L E  Microstructure evolution of a W-doped ZrB2 ceramic upon high-temperature oxidation  Laura Silvestroni1  | Diletta Sciti1  |  Fr\\x13ed\\x13eric Monteverde1  | Kerstin Stricker2  |  Hans-Joachim Kleebe2  1CNR-ISTEC, National Research Council  of  Italy  Institute of Science and  Technology for Ceramics, Faenza,  Italy  2TU-IAG, Technische Universit€at  Darmstadt   Institut  f€ur Angewandte  Geowissenschaften, Darmstadt, Germany  Correspondence  Laura Silvestroni, CNR-ISTEC, National  Research Council of  Italy  Institute of  Science and Technology for Ceramics,  Faenza,  Italy.  Email:  laura.silvestroni@istec.cnr.it  Funding information  European Community’s Seventh Frame work Programme, Grant/Award Number:  LIGHT-TPS No. 607182.  Abstract  This basic  research deals with the microstructure  evolution of  a W-doped ZrB2 ceramic, as-sintered and upon oxidation at 1650°C. Transmission electron micro scopy enabled to disclose microstructural  features occurred during oxidation never  observed before.  In the pristine material,  (Zr,W)B2  solid solutions  surround the  original ZrB2  nuclei, whereas  refractory W-compounds  at  triple  junctions  and  clean  grain  boundaries  are  distinctive  of  this  ceramic. After  oxidation,  the  microstructure is  typified by intragranular nanostructures,  in which nanosized W  inclusions  remained trapped within ZrO2 grains, or decorate  their  surfaces. The  understanding of  the oxidation reactions occurring in the system as a function of  the  oxygen  partial  pressure was  fundamental  to  conclude  that W-based  com pounds do not notably suppress or  retard the oxidation of ZrB2 ceramics.  K E Y W O R D S  ceramic,  inclusions, oxidation, SEM, TEM  1  |  I N T R O D U C T I O N  Zirconium diboride, ZrB2, compounds in the ultra-high temperature ceramics (UHTCs)  is one of  the most  investigated  class  of materials,  thanks  to  its  unique  combination  of  chemical and structural properties and its potential employ  in aerospace industry for  the development of  sharp hyper sonic vehicles or atmospheric most extreme conditions.1  entry probes  capable of  the  Some of  the recent  research activities in structural appli cations of ZrB2-based ceramics pointed out tion of WC, ranging from 5 to 20 vol%, provides benefits  that  the  addi not  only  to  the  oxidation  performances,  resulting  in  a  reduced thicknesses of  the scales  formed upon exposure to  high temperatures ics,2,3  for  both ZrB2 and HfB2 based ceramhigh-temperature strength,2-10 which experimental values exceeding 600 MPa were reported tests above 1500°C4,8 and more than 800 MPa at 1800°C.9 Some hypotheses have been suggested to explain  but  chiefly  to  for  the  role of W in the metal diboride matrices, but no con clusive mechanism has been proposed yet.  With respect  to the  effect of W compounds on oxida tion, it is reported that WO3 forms a eutectic with ZrO2 at 1273°C in the growing oxide scale during exposure to air at high temperature.11 This  liquid phase helps  the porous  outermost  scale to sinter oxygen.2  and thus  limits  the  inward diffu sion  of  Another  documented  phenomenon  involves  the  borate-glass  formed.  Kezmedashi  et al.  showed that 4 mol% W added to B2O3 stability at higher temperature, thus hindering its volatiliza increases  the glass  tion, so that  this factor could also be beneficial to the retenscale.12,13 The  tion  of  a  compact  outermost  oxide  same  conclusion was drawn upon experiments on HfB2 out at Wright Patterson Labs, where a more viscous phase  carried  separated glass was  found in the outermost  regions of  the  external oxide scale and the dense inner HfO2, slowing down the inward oxygen penetration to the sample core.3  Hypotheses  to  explain  the  excellent  high-temperature  flexure  strength  of  the ZrB2-SiC-WC system, which also be related to its oxidation behavior, were suggested by Zou et al.,4 who theorized that  can  strength retention, or  even  its  increase at high temperature, was due to oxide removal  Received: 3 October 2016  |  Accepted: 14 December 2016  DOI: 10.1111/jace.14738  1760  |  © 2017 The American Ceramic Society  wileyonlinelibrary.com/journal/jace  J Am Ceram Soc. 2017;100:1760-1772.  \\x0c', 'from grain boundaries and to a switch in the fracture mode  from transgranular  to intergranular.  In a  recent  study,  the  excellent mechanical  response of the ZrB2-SiC-WC system above 1500°C was explained  in an oxidizing environment  in terms of a combination of  favorable conditions,  includ ing the formation of a continuous glassy layer upon oxida tion that effectively healed surface flaws,  together with the  extensive formation of core-rim substructures with shell-toshell grain boundaries depleted of secondary phases.8,9 The  shells substructures, constituted by (Zr,W)B2 solid solutions with accumulation of dislocations at the core-rim interfaces,  were  considered  beneficial  for  triggering  a Hall  Petch  hardening  in  the  rims,  thus  enabling  plastic  deformation  to accommodate temperature.8,9  the  stresses  under  loading  at  high  Although all  these  explanations  are plausible, no basic  study has been undertaken to investigate the microstructure  evolution  at  nanoscale  level  upon  exposure  to  oxidizing  environment when W, or some of  its refractory compounds,  like WSi2, WB, and WC, are secondary phases of tered microstructure of a ZrB2 ceramic. Such investigations would be indeed fundamental to disclose the effective role  the sin of W in the improved performances of W-doped transition metal  diborides  in  view  of  tailoring  them  as  highly  performing structural ceramics.  In this respect, the oxidation behavior of a ZrB2-WSi2 material upon oxidation from 1200°C to 1800°C was previinvestigated:14  ously  it  was  observed  that  this  system  develops a complex multilayered scale containing zirconia,  a  partially  protective  silica-based  glass  and W-oxides.  In  the  subsurface  layer,  no  Si-depletion  layer  was  ever  observed  at any temperature, from 1500°C,  but  columnar ZrO2 convection phenomena  started  forming  owing  to  which activated the growth of zirconia in the form of pil lars, similar to what was observed for other ZrB2-SiC composites.15,16 The oxidation experiments at 1800°C showed  that no dense  and compact  external  layer or  stable oxida tion  products  in  the  subsurface  formed;  on  the  contrary,  pronounced migration of silica species and W-oxides to the  top surface induced the formation of owing to extensive gas escape.14  craters  and fissures,  In this paper, we describe upon oxidation at 1650°C of  the microstructure  evolution  a ZrB2 with 15 vol% of WSi2. Transmission electron microscope combined with chemical analysis were the added analytical  ceramic hot-pressed  tools here employed to disclose the chemistry, crystallinity,  and composition of  the new phases formed.  2  |  E X P E R IM E N T A L P R O C E D U R E  15 vol% WSi2 was added to ZrB2 to promote its densification and the following commercial powders were used:  hexagonal ZrB2 (grade B, H.C. Starck, Goslar, Germany), specific surface area 1.0 m2/g, particle size range 0.1-8 lm, impurity max content (wt%): C: 0.25, O: 2, N: 0.25, Fe:  0.1, Hf: 0.2,;  tetragonal WSi2 -325 mesh, 99.5%,  (Sigma Aldrich, Milwaukee, traces of metals <6000 ppm.  WI, USA),  The powders were weighed and mechanically mixed for  24 hours  in  absolute  ethanol  using ZrO2 milling media. dried using a rotary evaporator and 250 lm screen. A green diameter, was shaped by uniaxial pressing  The  slurry was  then  finally  sieved  through  pellet,  30 mm in  at  20 MPa. The pellet was then directly placed in the furnace 1930°C under  and  hot-pressed  up  to  vacuum (100 Pa)  using an induction-heated graphite die (32 mm inner diam eter) with an uniaxial pressure of 30 MPa 40 MPa and held at 1930°C for 40 minutes.  increased  to  The resistance to oxidation was studied using a bottom loading  furnace with MoSi2 FC18, Faenza, Italy) exposing 13.0 9 2.5 9 2.0 mm3 for 15 minutes to the effect of stagnant air at 1650°C. The coupon, previously cleaned in  heating  elements  (Nannetti  a  rectangular  coupon  acetone  and weighed, was  placed  onto  a  porous  zirconia  support, previously carved to limit  the contact between the  two surfaces,  so that  the coupon touched the zirconia sup port only at  the edges. The ceramic specimen was bottom up loaded into the hot zone of  the furnace when the target  temperature was achieved, elapsed 15 minutes, and rapidly  down-loaded from the hot chamber and let quench in air.  The microstructure before  and after oxidation was ana lyzed  on  fractured  and  polished  cross  sections  by  field emission scanning electron microscopy (FESEM, mod. ΣIGMA, ZEISS NTS Gmbh, Oberkochen, Germany) pled to an energy-dispersive X-ray microanalyzer (EDS,  cou mod.  INCA Energy  300, Oxford  Instruments, Abingdon on-Thames, UK). Key microstructural  features like residual  porosity, mean  grain  size,  and  volumetric  content  of  the  secondary  phases were  evaluated  from FESEM micro graphs elaborated with the support of  the commercial  soft ware  package  Image  Pro  Plus  (v.7, Media Cybernetics,  Rockville, MD, USA).  The samples for TEM analyses were prepared by cutting  a disk from the as-sintered pellet and from the coupon oxidized at 1650°C. The discs were mechanically ground down 20 lm and incipient perforations were observed by optical microscope.  to  about  then  further  ion  beam thinned  until  Local phase  analysis was performed using a  transmission  electron microscope  (TEM, mod.  JEM 2100F JEOL, Ltd.,  Tokyo, Japan) operating at a nominal voltage of 200 kV and  equipped with  an  energy-dispersive X-ray  system (EDS,  mod. INCA Energy 300, Oxford instruments, UK). Electron diffraction patterns identification was carried out through the software tool developed for Digital Micrograph.17  Phase  stability diagrams  and the  concentrations  of  the  most  relevant  compounds  in  well-defined  equilibrium  SILVESTRONI ET AL.  |  1761  \\x0c', '1762  |  conditions were  computed  by means  of  the  commercial  package HSC Chemistry v. 6.12 (Outokumpu research Oy,  Pori, Finland).  3  |  3.1  |  R E S U L T S A N D D I S C U S S I O N  As-sintered microstructure  Key microstructural  features of  the  as-sintered ceramic  SILVESTRONI ET AL.  sintered with  transition-metal  are depicted in Figure 1.  the matrix, with mean grain size  In analogy with other diborides disilicides or carbides,4-7,18 in the order of 3.5 lm, “core-shell”: pure ZrB2 and an epitaxial (Zr,W)B2 solid solution the shell, with a Zr-to-W atomic ratio of 98-to-2, Fig substructure defined as  constitutes  displays  core  the  a  ure 1A. The  estimated  volume  amount  of  solid  solution,  The crystalline phases detected by XRD in the as-sintered  close to 50% of  the diboride matrix, confirms an extensive  tetragonal WB with  development of  the W-doped shell.  Importantly, dislocation  ceramic were  hexagonal  ZrB2 amount. Actually,  and  WSi2 peaks  in minor  ZrB2 for 2-theta angles above 90 degree was attributed to  splitting  the  of  pile-ups  are present  at  the  core/shell  interface,  they could  be misfit  dislocations  related  to  a  small  lattice  parameter  the formation of a hexagonal  cell  parameters  if  shorter ZrB2.5  (Zr,W)B2 compared to  solid solution with  the  stoichiometric  mismatch between pure ZrB2 solution,18 Figure 1B. This result matches  the  conclusions  recently published  by Dai  rather well with et al.19  on the  and the W-containing solid  (A)  (B)  (C)  (E)  (D)  (F)  F I G U R E 1  FESEM and TEM images showing the microstructure of the as-sintered composite. (A) General overview evidencing the core-shell  structure of the matrix magnified in (B), (C) the secondary phases, (D) residual WSi2 trapped at the grain boundaries, (E) the formation of tetragonal  WB with the diffraction pattern inset, and (F) the structure of WB with the Fast Fourier Transformed image inset showing nanodomains  \\x0c', 'effects  of W on  the  shear  properties  of  (Zr,W)B2 calculations. The obtainment  deter mined  by  first  principle  of  specific  alloying,  (Zr,W)B2 the W-B bonds, compared to Zr-B ones,  in  the  present  case, weakens  reducing the stress  and the activation energy for nucleation of dislocations. agreement with,19  It  follows  that,  in  the  introduction  of W  into  the  ZrB2 at high  lattice  is  beneficial  to  the  retention  of  strength  temperature.  The  shells  represent  the  largely  distributed  reservoir  of W in  the  as-sintered  ZrB2-based ceramic object of In addition, W-containing phases,  the present work.  estimated via  image  analysis in a cumulative amount around 5 vol% thanks to a  bright Z-contrast in Figure 1A, were identified by TEMEDS as WB, WSi2 and W-C-O, as 3-5 lm large agglomerates or trapped as 600 nm triple junctions, Figure 1A,C,D,  E. All  these secondary phases contain traces of Zr.  In addi tion, WB,  a  reaction  product  formed  upon  interaction  between WSi2 and B2O3 shows 2-5 nm wide nanodomains,  covering  the  boride  particles,  Figure 1F,  indicating  possible ordered vacancies. About 3 vol% of SiO2 pockets trapping ZrO2 particles are also observed (Figure 1C). High-resolution TEM analysis in Figure S1 (supplementary  material)  reveals clean interfaces both in case of simple tri ple  junctions, Figure S1A,  and  in  case  of  segregation  of  secondary phases  at  the  triple points, Figure S1B. Several  interfaces were  analyzed  in  high-resolution mode  and  no  amorphous  film was  found  between  (Zr,W)B2, W-C-O, WSi2, and WB grains, as the examples in Figure S1C,D.  3.2 | Microstructure upon oxidation at 1650°C  XRD patterns 1650°C  from the  surface of  the  coupon oxidized at  revealed monoclinic  ZrO2 phase, with only little traces of the tetragonal polytype. Nei as main  crystalline  ther W-oxides,  like WO2 or WO3, nor WB were indexed. the oxidation attack induced a wavy and  On the  surface,  irregular  surface,  showing  craters  and  cabbage-like  struc tures, Figure S2A,B. The ZrO2 grains, 2 lm, are embedded into a continuous EDS spectra in Figure S2A. Surface  as  coarse  as  about  silica layer,  see the  uplifts  and  growing  planes  associated  to  the tetragonal-to-monoclinic phase transformation of ZrO2 20 are well visible in Figure S2C. FESEM images of  the oxidized fractured cross  section  in Figure 2 show a multilayered architecture. Three princi pal  subscales  and two interfaces  are  identified and tagged  in Figure 2A. A discontinuous  surface silica scale,  labeled  S, a few micrometers scale 40 lm thick, Figure 2B. No continuous  thick,  covers  a mixed  SiO2/ZrO2 labeled ZS in Figure 2A and zoomed in  external SiO2 Si-carrying phases  skin formed due  to  the  lack  of  residual  after  sintering  and to the coarsening of ZrO2 grains which made the surface very uneven.  (A)  (B)  (C)  (D)  (E)  F I G U R E 2  SEM images of  the  cross  section after oxidation at  1650°C. (A) Overall appearance with outline of the magnified areas in (B-E). S: SiO2, ZS: ZrO2 + SiO2, ZW: ZrO2 + W-compounds, Bulk:  un-oxidized ceramic.  (B) Outermost coarse ZrO2 with silica layer on  top,  (C)  interface between the outermost ZrO2 and the columnar ZrO2  scale,  (D) W-oxide phase  and W-nanobeads  spread across  the ZrO2  scale  and (E)  sharp interface between the oxide  and the boride un oxidized bulk  SILVESTRONI ET AL.  |  1763  \\x0c', 'The second layer  is a 180 lm columnar ZrO2 partially labeled ZW in Figure 2A. The columnar  filled with silica,  grains of  this ZW layer  exhibit  enhanced directionality in  the upper part of this subscale. In proximity of the face 1 shown in Figure 2C, W-oxides, as large as 3 lm are found among ZrO2 grains and their concentration readily decreases toward the un-oxidized bulk. A magnified image  Inter of  such W-oxides  is  shown  in Figure 2D.  In  addition,  a  dense  population  of  ultrafine  particles  characterized  by  bright  Z-contrast  and  spherical  shape,  some  tenths  of  nanometers in diameter, and hereafter named as nanobeads,  can  be  seen  in Figure 2D. Such  nanobeads  are  homoge neously dispersed throughout  the  entire  columnar  zirconia  ZW sublayer.  In general,  the deepest bottom of  the oxide  layer,  i.e.,  Interface 2,  is well anchored to the unoxidized bulk with a  sharp  interface,  Figure 2E,  suggesting  that  the  oxidation  front  proceeds  transgranularly  instead  of  intergranularly,  similar to what was reported for other ZrB2-based composites containing W-based compounds.8  TEM-EDS analyses of  the oxidized sample  enabled to  disclose finer microstructural details and will be discussed  in the  following with reference  to the  layers  identified in  Figure 2A.  S:  First,  the  outermost  thin  coating was  found  to  be  made  of  amorphous  silica, without  traces  of  crystalline  nanoprecipitates,  see Figure 3. The  features with different  contrast  resolved in Figure 3 are chemical variations of  the  SiO2-based glass that Zr, W, and other impurity cations,  induced phase separation and include  like Ca and Al.  ZS: The outermost ZrO2 grains series of defective structures like  in this  layer  exhibit  a  dislocations,  twinning,  and stacking faults, as a result of  the tetragonal-to-monocli nic  transformation upon cooling, Figure 4. No trace of W  was detected by TEM-EDS analysis inside the ZrO2 grains, confirming the negligible solubility of W in the zirconia  lattice. The glassy phase which permeates  the boundaries  of  the  columnar ZrO2  grains  has  a  varying  composition  containing traces of Zr, W, and Al, but,  in contrast  to the  outermost one (see Figure 3), crystalline precipitates can be  easily found, see the rounded area in Figure 5.  Interface 1:  In this region,  faceted 200-400 nm particles  with dark contrast and hexagonal  shape are found embed ded in a SiO2-based analysis coupled to  glassy phase, Figure 6A. TEM-EDS  electron  diffraction  revealed  these  faceted particles  to be  cubic W with Zr  in a  ratio around  85:15 at.%,  Figure 6B.  The  glassy  phase well wets  the  metallic  particles, Figure 6B-D. Also  at  this  depth, W is  not  found in monoclinic ZrO2 grains, as confirmed by the TEM-EDS spectra in Figure 6C. However, in proximity of  these W,Zr metallic  particles, monoclinic WO3 whiskers are often found, Figure 7. This phase has a fibrous aspect  when  not  oriented, Figure 7A and  assumes  a whisker  or  rod-like  shape  once  oriented, Figure 7B,D. The  electron diffraction image of  the whisker  in Figure 7C shows  that  the growth direction of  the monoclinic WO3 occurred along and a striking feature is found in the  the  [001] direction  diffraction  spots. This  is  due  to  stacking  faults  possibly  related to oxygen deficiencies, that are commonly observed in tungsten oxides.20,21 The WO3 SiO2-based glass, containing traces ure 7E.  rods  are  surrounded by  of W and  Zr,  Fig These  phases  could  be  the  oxidation  products  of  the  secondary phases present  in the as-sintered microstructure.  This assumption is based on two facts. First,  these pockets  have  similar  size  as  the  pristine W-based  phases  (few  micrometers)  and secondly,  the metallic particles  contains  comparable amount of Zr, as detected in the residual WB,  WSi2, and W-C-O. ZW layer: The TEM-EDS analyses  of  this  oxide  sub scale  confirmed the  characteristic  columnar growth of  the  monoclinic ZrO2 with nanobeads of variable size from 5 to 60 nm, Figure 8.  grains  externally  and  internally  adorned  Those nestled inside ZrO2 grains the advancement of the dislocations  sometimes appear  to pin  front,  Figure 8C.  High-resolution  TEM images  coupled  to  structural  and  F I G U R E 3  S layer, TEM image of  the oxidized composite showing the morphology and composition of  the outermost glass among mono clinic ZrO2 grains with corresponding diffraction pattern and EDS spectra  1764  |  SILVESTRONI ET AL.  \\x0c', 'EDS analyses  revealed these nanobeads  to be metallic W  with  some  oxygen  trace, Figure 9. Their  spherical  shape  suggests  a very low wettability of W toward zirconia.  In  addition,  half-moon  shaped  residual  glassy  pockets,  as  shown  in  the  high-resolution TEM image  of  Figure 9C,  were systematically found adjacent  to such nanobeads.  The morphology  and  arrangement  of  such  zirconia  grains with  encapsulated W inclusions  presented  in  Fig ure 9 let us deduce that  these are the oxidation products of  the original  (Zr,W)B2 solid solution, as it will be discussed later. The W nanobeads remained un-oxidized upon their  formation:  this  finding  leads  the  authors  to  infer  that,  within the  zirconia scale marked as ZW in Figure 2A,  the  oxygen partial pressure is  low enough to leave metallic W  as most  favorable condensed phase. On the contrary, it was at 1650°C, W released in proximity of  deduced that,  the  external surface volatilize,22  tended to first  form liquid WO3 inclusions-free ZrO2 grains progressively filled the voids upon coalescence  and then  thus  leaving  that  and coars ening mechanisms at high temperature.  3.3  |  Thermodynamic considerations  To  better  correlate  the microstructure  evolution  to  the  oxidation  response  of  this W-doped ZrB2  ceramic,  some  thermochemistry  considerations  were  addressed  through  the HSC Chemistry commercial  software package. Unfor tunately,  the  (Zr,W)B2 phase commercial package,  is not  included in the data base  of  the  so  we  could  not  directly  compare  its  tendency  to  oxidize with  that  of  pure ZrB2. Due this paragraph we  to  this  lack  of  thermodynamic  data,  in  can  track  the  oxidation  of  the  pure  ZrB2 WB, WSi2, involving the  compared  to  pure metallic W and  its  compounds,  and WC,  whilst  the  chemical  reactions  (Zr,W)B2 a qualitative  solid  solution with  oxygen will  be  treated  in  approach  and  supposed  to  have  an  oxidation  behavior  intermediate  between  that  of  ZrB2 A  and WB.  number  of  thermodynamically  possible  oxidation  reactions  involving ZrB2 WB and WC are listed in Table 1. However,  or  the  secondary  phases WSi2, for the oxida tion mechanisms more favorable in the subsurface scales,  it  is more  illustrative  to look at  the phase stability diagrams  in Figure 10. These maps are rather useful  to define stabil ity areas of condensed phases vs temperature and/or chemi cal potential  for selected systems.  The plot  in Figure 10A shows  that  at 1650°C, marked  by a dotted line,  the oxygen partial pressure  (PO2) has role in the nature of condensed phases that we  a  fundamental  find in the microstructure: W is  stable  for oxygen partial  (A)  (B)  (C)  F I G U R E 4  ZS layer, TEM images, and diffraction patterns of  the HRTEM pictures in the boxed areas evidencing defects in the outermost  ZrO2 grains pointed by arrows:  (A) dislocations,  (B)  twinning,  (C) stacking faults. The EDS spectrum confirms no W trace within ZrO2 grains  SILVESTRONI ET AL.  |  1765  \\x0c', 'pressure  below ~10-7 atm, while  above  this  value  liquid  WO3 in the form of nanobeads all  is more favorable. This explains why W can survive  throughout  the subsurface and  why WO3(s) was columnar ZrO2 scale, Solid WO2 was never due to its narrow stability domain.  found mostly in the upper  zone of  the  i.e., at higher oxygen partial pressure.  found in the oxidized microstructure,  To simplify our  thermodynamic  computations, we con sidered  two  pressure  conditions:  1 atm, representative and 10-7 atm,  of  events occurring on the  surface,  to account  for  the decreased pressure in the inner portions of  the mul tilayered oxide scale* . Then,  to get an insight  into the oxi1650°C,  dation  effects  on  the microstructure  evolution at  the thermodynamics of ZrB2, WB, WSi2, WC, and W were addressed by equilibrating 1 mole of each solid component with 2.5† mole of oxygen. The main outcomes of putation through HSC Chemistry can be summarized  the com as  follows.  What happens on the surface is straightforward (1 atm):  of  the initial ZrB2 we find just ZrO2 and liquid B2O3. As  for  the secondary phases, WB produces a mixture of B2O3 and WO3 in glassy form plus complex gaseous W-O oxides like  (WO3)x, WSi2 while WC mainly gaseous  yields metallic W plus  SiO(g),  species. The W developed from  the  oxidation  of W-boride,  silicide,  and  carbide  has  no  chance to survive and evaporate agreement with.22  as complex W-oxides,  in  In conditions of  low pressure (10-7 atm), ZrO2 now in vapor phase. Similarly, all  remains  and B2O3 compounds  is  the W develop  large  amounts  of  corresponding  gas eous  species.  From these  outputs, we  argue  that  in  the  subsurface layers  the partial pressure of  the various gases,  like BxOy, WxOy, local chemistry of  SiO,  and COx, condensed oxidation  has  influence  on  the  products,  with  B2O3(g) playing the major In this context, to get an  role, due  to the boride matrix.  insight  into  the  hierarchy  of  oxidation reactions of borides, Figure 10B,C presents sep1650°C for  arate  phase  stability  diagrams  calculated  at  the Zr-B-O and W-B-O systems  varying  the  oxygen  and  boron  oxide  partial  pressure. Looking  at  the  highlighted  areas  in  Figure 10B-C, we can read that if we set fix temperature, 1650°C, and oxygen and boron  conditions of  oxide  partial  pressures,  the  stable  condensed  phase  are  ZrO2 firming  in one  case, but W plus WB in the other one  con our  experimental observations. This  also suggests  that ZrO2 des, so that  is  thermodynamically more  stable  than W-oxi the  limited  amount  of  oxygen  diffusing  into  the  oxide  is  preferentially  gettered  by  the Zr-containing  phases,  rather  than by W.  3.4  | Oxidation behavior  The microstructural  features observed in the oxidized speci men, supported by thermochemical equilibrium predictions,  were  used  to  formulate  hypotheses  over  the  oxidation  mechanisms which ceramic at 1650°C.  involved  the  present W-doped  ZrB2  The matrix of  the as-sintered material,  sketched in Fig ure 11A is constituted by ZrB2 grains homogeneously surrounded by about 50 vol% of (Zr,W)B2 solid solution, crystalline W-containing  then  around  5 vol% of  phases  (WB, WSi2, tion) are segregated  and WC in  decreasing  order  of  concentra at  the  triple  junctions. Also,  about  3 vol% of silica is present, WSi2 during sintering.5 The oxidation of ZrB2 at 1 atm yields liquid B2O3 and solid ZrO2. The solid solution around the ZrB2 grains oxidizes as well: given the negligible solubility of W in ZrO2, diagram11 as reported from the Zr-W-O phase and con as  dissociation  product  of  firmed by TEM-EDS analyses  in Figures 4C, 6C and 9A,  W is released as metal within the growing ZrO2. The overall oxidation of the diboride matrix can be expressed as  follow (11):  F I G U R E 5  ZS layer, HR-TEM image of  the intergranular glass  running  through  the  columnar  ZrO2  grains,  showing  traces  of  crystalline  compounds  in  the  SiO2-based  glass  (circled  area). The  insets are: a low magnification image of  the zone,  the EDS spectrum,  and the diffraction pattern of the partially crystallized glass  *This pressure  level was  arbitrary selected as  representative of  the  subsurface  scale by looking at  the coexistence field of W and WO3 in Figure 10A.  †The  amount of oxygen was chosen to allow the  completion of  the oxidation  reactions of ZrB2, see Table 1.  1766  |  SILVESTRONI ET AL.  \\x0c', 'Zr1\\x00xWx  ð  ÞB2 þ 2:5 \\x00 x  ð  ÞO2 ! 1 \\x00 x  ð  ÞZrO2 þ xW þ B2O3 11  ð  Þ  Of  the solid products, only ZrO2 was found in the outer surface (ZS in Figure 2A) since W transforms into WO3(l) and then volatilizes.23-28  The  voids  left  by  the  progressive  oxidation  and  volatilization of  the metal  inclusions  from the  surface  are  filled by the growing ZrO2, explaining a multistep process of cell distortion and defects creation, Figure 4. On the  contrary,  in  the  subsurface  scale  at  low  pressure, W  encased in ZrO2 grains tion of the W nanobeads seems localized in ZrO2 domains, the oxide preserves the former core-shell morphology.  remains un-oxidized. The distribu i.e.,  The characteristic droplet-like shape might suggest  that  liq uid W experienced a fast quenching and, driven by reduced  wettability toward ZrO2, all the ZrO2 volume, Figure 11B-C. At the same time, residual WSi2, WB, and WC have no chances to survive un-oxidized and start yielding W, liquid  took a globular  shape throughout  SiO2, and gaseous species like CO, SiO, B2O3, and W oxides. The availability of silica and borosilicate glass provides  prompt protection against  the initial oxygen diffusion. Fur ther  contribution to the  formation of  the outermost glassy  coating derives from the SiO2, already present in the material  upon sintering, and B2O3 coming also from the oxidation of the diboride matrix. B2O3 dissolved into silica decreases the viscosity of the in situ formed boro-silicate glass.29 As soon  as an outermost  local silica lid is formed,  it acts as localized  protection for  the products underneath,  therefore W-oxide  and metal residues can be retained as solid compounds, Fig ure 11C. The  local  formation  of W(Zr)  particles  in  this  region in hexagonal shape could be due to variations in the  vapor conditions, which strongly influence morphology of the condensed products.30  the  shape  and  As long as the oxidation front advances deeper  into the  bulk,  the oxide  scale  further  thickens,  thus  extending the  diffusion  distance  for  oxygen  to  reach  the  oxide/boride  interface.  It  follows  that  the oxygen partial pressure of  the 10 regions overlying 7 atm established  the  interface  does  not  exceed  the  by the  thermochemical  equilibrium pre dictions:  under  these  conditions, W is  safe  inside/outside  the ZrO2 grains.  The  role of other oxidation gaseous products  like CO  (g), SiO(g), and volatile boron oxides sidered: at 1650°C,  should be also con they tend to evolve and escape outside  the  oxide  scale,  Figure 11D, wherever  they  are  formed  and/or captured in some glassy melts.  (A)  (B)  (C)  (D)  F I G U R E 6  Interface 1, (A) TEM image of W,Zr particles immersed in the glass within the columnar ZrO2 grains with low-magnified view of  the area inset. The HR-TEM images in (B-D) are magnified images of the boxed areas in (A) and show the composition of ZrO2, particle and glass  SILVESTRONI ET AL.  |  1767  \\x0c', '1768  |  (A)  (B)  (D)  SILVESTRONI ET AL.  (C)  (E)  F I G U R E 7  Interface 1,  (A),  (B), and (D) TEM images  showing the formation of WO3 whiskers with diffraction patterns.  (C) and (E) are  magnification of  the boxed areas  in (B) and (D),  respectively with EDS spectra  (A)  (B)  (C)  F I G U R E 8  ZW layer,  (A) and (B), TEM images of  typical columnar ZrO2 grains in ZrB2-WSi2 composite after oxidation showing encased  W-nanobeads. Note in (C)  the dislocation advancement arrested by the metal spheres  \\x0c', 'More  generally,  the  pronounced  directionality  of  the  columnar  zirconia grains  is  commonly ascribed to release by-products.31 Actually,  of  the  aforementioned  gaseous  sublayers made of columnar zirconia were observed also in W-free oxidized ZrB2 ceramics,1,3,15,16,31 i.e., of the release of volatile W oxides. It follows  independently  that  the cre ation of such configuration cannot be directly linked to the  formation, evolution and outward escape through the oxide  scale of  the only volatile W oxides. At  the same time,  it  is  logically obvious to link the oriented grain growth to exten sive diffusive mechanisms along the grain boundaries and/or  channels which interconnect the zirconia pillars.  Verified  that W survives  un-oxidized where  zirconia  exhibits  the  columnar  shape  and that  the  channels  formed  among its boundaries are occupied by a glass,  the vigorous  outward diffusion of fluid by-products  is proposed as  the  driving force  for  the oriented growth of  the  zirconia  sub scale. At  the roof of  the columnar scale,  i.e., at  Interface 1  in Figure 2A, SiO(g) and volatile boron oxides  recondense  and  replenish  the  glassy  component  of  the  outer mixed  SiO2/ZrO2 scale:  in such a melt,  the initially formed nuclei  (A)  (B)  (C)  F I G U R E 9  ZW layer,  (A) TEM image of W nanobeads nestled in pure ZrO2 with EDS spectra and diffraction patterns and HR-TEM images  in (B) and (C). Note in (A) and (C)  the presence of glass adjacent  to the metal  inclusions  T A B L E 1  List of possible reactions between the major solid phases  of the as-sintered ceramic with oxygen  N°  Reaction  1  ZrB2 + 2.5O2 (g) ? ZrO2 + B2O3(l)  2  WB + 2.25O2 (g) ? WO3(l) + 0.5 B2O3(l)  3  2WB + 3.5O2 (g) ? 2WO2 + B2O3(l)  4  2WB + 1.5O2 (g) ? 2W + B2O3(l)  5  WSi2 + 3.5O2 (g) ? WO3(l) + 2SiO2(l)  6  WSi2 + 2O2 (g) ? W + 2SiO2(l)  7  WSi2 + 2.5 O2 (g) ? WO3(l) + 2SiO(g)  8  WSi2 + O2 (g) ? W + 2SiO(g)  9  WC + 2O2 (g) ? WO3(l) + CO(g)  10  WC + 0.5O2 (g)? W + CO(g)  SILVESTRONI ET AL.  |  1769  \\x0c', '1770  |  (A)  (B)  (C)  SILVESTRONI ET AL.  F I G U R E 1 0  (A) Phase stability diagram of the W-O system as a function of oxygen partial pressure, PO2(g), and temperature, T. The dotted line marks the experimental temperature of 1650°C. Phase stability diagrams of condensed phases as a function of oxygen and boron oxide partial pressures calculated at 1650°C for  in the same conditions of  the most stable compound is different  the (B) Zr-B-O and (C) W-B-O system. Gray areas in (B) and (C) evidence that, in the two systems, ZrO2 vs W + WB  temperature and pressure,  (A)  (B)  (C)  (D)  F I G U R E 1 1  Sketch of  the oxidation phenomena occurring at 1650°C in the W-doped ZrB2 composite [Color figure can be viewed at  wileyonlinelibrary.com]  of zirconia find favorable conditions  to coarsen due to the  presence of  the glass, Figure 11(D).  Therefore,  according  to  the microstructure  evolution  here  presented  and  to  previous analysis on the oxidation temperatures up to 1800°C,14  behavior of  this composite at  the resistance to oxidation of ZrB2 ceramics does not seem to take great advantage solely from the presence of W temperature reaches or exceeds 1650°C:  when the exterior  the progressive volatilization of W-oxides  from the surface  does not offer any further protection to the material  integ 3.5 | What is the role of W in the oxidation behavior of ZrB2?  In the  light  of  the  new microstructural  features  disclosed  by  TEM analysis, we  addressed which  role W and  its  refractory  compounds, WB, WC,  and WSi2, oxidation behavior of a ZrB2 ceramic. Our attention is primarily devoted to the (Zr,W)B2 solid solution which is the widest W-based compound in the as-sintered ceramic.  play  the  in  First, W has an indirect  impact, because, upon sintering,  rity. On the other hand,  the  authors  are  convinced of  the  it promotes the development of a microstructure with clean  major  effectiveness  coming  from the  Si-carrying  phases,  grain  boundaries  and  only  refractory  phases  at  the  triple  residual WSi2 improving the  and  SiO2 resistance  glassy pockets in this case, in to oxidation up to 1650°C when  added to ZrB2 matrix.  junctions. As  a  consequence,  the oxidation front proceeds  coherently, attacking the diboride matrix in a transgranular  mode,  instead of an intergranular one. This  is mainly due  \\x0c', 'to the absence of  soft phases at  the triple points, and also  to the lower  tendency of  the WB phase to oxidize, as com pared to pure ZrB2 tion boundary can be beneficial to 1650°C, but we can expect  (see Figure 10B,C). This  even oxida for  temperature regimes up  that an excessive coarsening  of the columnar ZrO2 observed in.14  layer will  lead  to  spallation,  like  As for  the direct  role of W on the response to oxidation  of  the ZrB2 able in the  ceramic,  two cases  are pointed out: W avail solid  solution,  or  constituting  the  secondary  phases WB, WSi2, and WC. In the case of the (Zr,W)B2 solid solution, released W turns into finely dispersed W nano it was shown  that  the  beads encapsulated in ZrO2 grains, or decorating their faces. When the oxygen partial pressure overcomes 10-7 atm,  sur these metallic nanobeads  convert  into liquid W oxides. That  reaction involves  a volume  expansion which  might partially contribute to close channels among zirconia  grains,  retarding  the  inward  ingress  of  oxygen  accessing  through  them.  For  oxidizing  temperature  at which  the  volatilization  of  the W oxides  is  suppressed,  this mecha nism can be more effective. However,  the formation of oxi des  that,  at  certain  depths  of  the  ceramic,  are  stable  as  condensed phases or cations  (Zr, W)  that dissolved in the  glass increase its viscosity and induce precipitation of crys talline phases (see Figure 5), could be favorable conditions  to retard oxidation.  Regarding the formation of an eutectic between ZrO2WO3,11 proposed as mechanism able to retard the oxidation,2 the present investigations did not  reveal any trace of  W in the final  crystalline ZrO2, but we exclude the formation of such liquid protective phase.  cannot ultimately  In addition,  some  authors  suggested that W oxide  spe cies form strong acid sites in ZrO2 and inhibit ZrO2 tetragonal to monoclinic polytypic transformations.32 The  findings of  the present work supported by X-ray and elec tron  diffraction  revealed  only  the monoclinic  polytype  at  any depth of  the oxide scale: studies,5,14  in this  respect  the  authors,  according  to  other  are  inclined  to  attribute  a  minor  role to W in the stabilization of  tetragonal ZrO2. secondary phases distributed throughout  As  for  the  the  diboride matrix, WB, WSi2, and WC, cated as the most effective to improve the resistance to oxi the disilicide is indi dation because of the delivery of silica glass which is known  to be an excellent sealing agent as well as a barrier against temperature up to 1650°C. On the other  oxygen diffusion at  hand, WB,  taken as comparison term associable to the (Zr,  W)B2 ZrB2, Figure 10B,C. WC oxidizes to W or volatile W-oxides (reactions 9,10) and additional WB phase could further  solid  solution,  has  lower  tendency  to  oxidize  than  locally form by recombination of W and boron oxide.  Borides  like ZrB2 and its boron oxide, contributing to accelerate the high-temperature  solid solutions provide liquid  oxygen diffusivity through the less viscous glass, compared to a purely formed silica,29  since  the domain stability  of  the in situ formed boro-silicate glass widens upon capturing  part of  the available W/WO3, as proved in dedicated stud ies.12-14  The  authors’ final of W by alloying, i.e., dissolved into the ZrB2 strengthening ZrB2 ceramics core-shell substructures,9  speculation is  that  the  incorporation  lattice,  is  mostly  effective  in  through  the  development  of  rather  than  retarding  or  suppressing  the  overall  inward  diffusion  of  oxygen.  4  |  C O N C L U S I O N S  A ZrB2 thanks to the addition of 15 vol% WSi2, was analyzed by transmission electron microscopy in the as-sintered state 1650°C in  ceramic,  hot-pressed  at  1930°C to  full  density  and  upon  oxidation  at  static  air. The  original  matrix was composed of ZrB2 cores B2 solid solution shells, with WB, WSi2, tory phases at the triple point junctions,  surrounded by (Zr,W)  and WC refrac all  characterized  by grain boundaries free of amorphous phases.  This  ceramic  after  high-temperature  oxidation  evolved  into a multilayered architecture with a thin outermost SiO2based glass, a coarse rounded ZrO2 scale partially filled with glass, a columnar ZrO2 well anchored to the unreacted bulk. TEM observations coupled to EDS analyses con firmed a negligible  solubility of W into the ZrO2 formation of W nanobeads encased  lattice,  but  rather  the  into  columnar ZrO2 grains or decorating their surfaces. W particles and WO3 whiskers embedded into silica-based pockets were also identified across the oxide scale at different  depths, depending on the local oxygen partial pressure.  Detailed microstructural  analysis  combined with  the  thermodynamics 1650°C or  equilibria  allowed  to  propose  that  at  above,  the main  role  of W during  the  oxida tion  of ZrB2 thus it  is  related  to  the  release  of  volatile W oxi des,  does  not  provide  notable  protection  to  the  advancement of  the oxidation front. The presence of  little  amount  of WB, which  is  less  prone  to  oxidize  than  Zr-boride,  and WSi2, tion-protective boro-silicate  enabling the  formation of  an oxida glass,  partly  mitigated  the  oxidation attack. at 1650°C in air,  Therefore,  the  formation of  the  core shell  configuration seems more  effective  in improving the  high-temperature flexure  strength of  a ZrB2  ceramic,  than  protecting it  from oxidation.  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  "_id": 129,
  "PDF": "Microstructure, mechanical properties and oxidation behavior of TaC- and HfC-based materials containing short SiC fiber.pdf",
  "Text": "[\"Available online at www.sciencedirect.com  Ceramics International 41 (2015) 1367-1377  CERAMICS  INTERNATIONAL  www.elsevier.com/locate/ceramint  Microstructure, mechanical properties and oxidation behavior of TaCand HfC-based materials containing short SiC fiber  L. Pienti, L. Silvestronin, E. Landi, C. Melandri, D. Sciti  CNR-ISTEC,  Institute of Science and Technology for Ceramics, Via Granarolo 64,  I-48018 Faenza,  Italy  Received 5 June 2014;  received in revised form 1 September 2014; accepted 12 September 2014  Available online 19 September 2014  Abstract  TaCand HfC-based composites containing 15 vol% short SiC fibers as reinforcing phase were produced by hot pressing at 1700-1750 1C. Suitable sintering additives were selected in order to get a full densification at temperatures well-tolerated by SiC fibers. Microstructural  characterization outlined a strong interaction between the carbide matrix and SiC fibers. The following properties were evaluated: Vickers hardness, fracture toughness by chevron notched beams, elastic modulus, 4-pt bending strength at room temperature and at 1200 1C, linear CTE up to 1300 1C in Ar, and thermal conductivity up to 1900 1C in Ar. Oxidation tests in air were carried out with thermogravimetric analysis up to 1500 1C and in a bottom up furnace at 1600 1C. TaC-based materials generally possessed higher thermo-mechanical properties compared to HfC, but lower oxidation resistance. Properties of fiber-reinforced composites were compared to those of unreinforced ones.  & 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  Keywords: Fibers; TaC; HfC; Thermo-mechanical properties  1.  Introduction  Among advanced ceramics, tantalumand hafniumbased refractory carbides are promising materials for applications in the aerospace and energy sectors. These materials combine the physical properties of ceramic and the electronic properties of metals. They are stiff, with Young's modulus values competing with those of SiC, and have good thermal conductivity, permitting heat to be drawn away from the superheated surfaces. Furthermore, they have high melting points (4 3900 1C) and good chemical resistance. These properties make TaC and HfC interesting systems for extremely hightemperature applications such as rocket nozzles, scramjet components, and future generation thrusters that require resistance to severe combustion conditions [1-3]. Mono l i th ic TaC and HfC are hard to dens ify due to the ir s trong cova len t bonds and low se lf-d if fus ion coeffic ien ts [1-3] . Ho t press ing , spark p lasma s in ter ing and pressureless s in ter ing have been of ten assoc ia ted to the use o f  nCorresponding author. Fax: þ 39 0546 46381. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  http://dx.doi.org/10.1016/j.ceramint.2014.09.070  0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  add i t ives to imp rove dens ifica t ion and m i t iga te the severe s in ter ing tempe ra tures [4-21] . Sma l l amoun ts o f me ta ls such as Fe , Mn , Co , and N i , lower the tempera tu re needed for dens ifica t ion [5-7] ; a l terna t ive s in ter ing add i t ives such as B , C , and B4C , remove surface ox ides on the s ta r t ing 1900-2300 1C par t ic les enab l ing dens ifica t ion a t [8 ,9 ] . Me ta l d is i lic ides , such as TaS i2 and MoS i2 have shown to be effec t ive add i t ives a l low ing dens ifica t ion of TaC and HfC a t tempera tures be tween 1750 and 1950 1C [11] e i ther by ho t press ing or pressure less s in ter ing [21] . Reduc ing par t ic le s ize of the powders has a lso been shown to dens ifica t ion . TaC and HfC þ 5 vo l% MoS i2 improve the w ith re la t ive dens i t ies above 90% have been ob ta ined s tar t ing from syn thes ized u ltrafine powders [10 ] . Recen t ly , spark p lasma s in ter ing was a lso used to conso l ida te Ta0 .8H f0 .2C us ing MoS i2 and TaS i2 as s in ter ing a ids a t 1650 1C under 30 MPa [19] . Published values for the thermo-mechanical properties of these carbides are still scarce. Hardness of 21-27 GPa and fracture toughness of 3.5 MPa m1/2 were measured by Balani et al. on vacuum plasma sprayed TaC [22]. TaC Young's modulus value of 550 GPa was measured by Lopez-de-la-Torre et al. [23]. For hot  \\x0c\", \"1368  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  pressed HfC-based composites with 15 vol% MoSi2, hardness up to 19 GPa was obtained [14], and mechanical strength of 450- 465 MPa and Young's modulus of 415-434 GPa for pressureless composites with 5-10 vol% of the same sintering additive were measured [16]. There are no recent reports on thermal properties like linear CTE up to 1300 1C in Ar, thermal conductivity up to 1900 1C in Ar. One drawback of tantalum and hafnium carbides is the poor oxidation resistance [1-3]. TaC oxidizes to Ta2O5 which has a relatively low melting point, 1800 1C [24]. According to previous studies on the high temperature oxidation of HfC-TaC compo[25,26], at To 1800 1C the outermost oxide layer sites formed was constituted of Ta2O5, while at T4 1800 1C it was made of a mixture of Ta2Hf6O19 and HfO2. HfC oxidizes to monoclinic HfO2, one of the most stable refractory compounds with a melting point of 2800 1C [24]. Oxidation tests carried out between 1400 and 2060 1C [27] and arc jet tests between 2400 and 2700 1C [28] confirmed the excellent oxidation response in severe condition of monolithic HfC. Similarly, HfC composites with MoSi2 at 5 vol% showed a good thermal stability when tested in an arc jet facility up to 2400 1C [29]. In this work the effect of SiC fibers on microstructure, thermomechanical properties and oxidation resistance of TaCand HfCbased composites was analyzed. TaSi2 has been selected as a sintering additive for the hot pressing process as it allows densification at relatively low temperature, compatible with the stability limit of SiC fibers. Interaction between the carbide matrix and SiC fibers was studied. The following properties were measured: Vickers hardness, fracture toughness by chevron notched beams, elastic modulus, 4-pt bending strength at room temperature and at 1200 1C, linear CTE up to 1300 1C in Ar and thermal conductivity up to 1900 1C in Ar. The oxidation behavior was tested through thermogravimetric balance up to 1500 1C and in a bottom up loading furnace at 1600 1C.  2. Experimental procedure  produced  starting  from commercial  The composites were raw materials: \\x0f TaC (Treibacher Industrie AG, Althofen, Austria), mean particle size 1.1 μm; \\x0f HfC (Cerac Inc., Milwaukee, WI, USA), particle size range \\x0f TaSi2 (ABCR GmbH & Co., Karlsruhe, Germany), particle 0.2-1.5 mm; size o 45 mm;  \\x0f SiC fibers (Tyranno SA3, UBE Industries, Ltd.) diameter 7.5-10 mm, composition wt% Si:C ¼ 67:31, oxygen o 1 wt%, and aluminum o 2 wt%.  labeled as reported in Table 1,  The following compositions, were prepared (vol%): \\x0f TaC þ 10 TaSi2 þ 15 chopped SiC fibers (TCTf); \\x0f HfC þ 10 TaSi2 þ 15 chopped SiC fibers (HCTf).  For the sake of comparison, the respective matrices were e.g.: TaC þ 10 vol% TaSi2 also produced and characterized, (TCT) and HfC þ 10 vol% TaSi2 (HCT). The powders were dispersed in absolute ethanol by simultaneously applying magnetic stirring and ultrasonication. For the reinforced composites, fibers were in house chopped. Low magnification SEM-images of the chopped SiC fibers and of a fiber section are reported in Fig. 1. The powder mixtures were further ball milled in absolute ethanol for 24 h, at 130 rpm, with SiC milling media. A fixed powder/milling media/solvent weight amount (1:1:1) was adopted. After milling, the slurries were dried in a rotary evaporator, sieved and shaped in four centimeters diameter pellets. The pellets with fibers were debonded at 500 1C for 1 h and all the samples were hot 1700-1900 1C range pressed at temperatures in the in low vacuum (100 Pa) using an induction-heated graphite die with a uniaxial pressure of 30 MPa, that was increased to 40 MPa during the isothermal stage; free cooling followed. For each composition, the maximum sintering temperature was set on the basis of the shrinkage curve, Table 1. After sintering, the bulk densities were measured by Archimedes' method. Crystalline phases were identified by x-ray diffraction (Bruker D8 Advance, Bruker, Karlsruhe, Germany). The microstructure was analyzed using a scanning electron microscope (FE-SEM, Carl Zeiss Sigma NTS GmbH, Oberkochen, Germany) and an energy-dispersive spectroscopy (EDS, X-Act, INCA Energy 300, Oxford Instruments, Abingdon, UK) on fractured and polished surfaces. Sections of the sintered materials were cut perpendicularly to the hot pressing direction and polished with diamond paste to 0.25 mm. Mean matrix grain size, amount of porosity and amount of secondary phases were determined on micrographs of polished sections using image analysis (Image-Pro Analyzer 7.0, Media Cybernetics, Silver Spring, MD, Rockville, USA). The mean grain size of the matrix grains was calculated by the circle  Table 1  List of  the samples:  label, sintering conditions (TMAX, dwell  time, pressure), density, matrix mean grain size and maximum fiber  length.  Label  Sintering (1C, min, MPa)  Th. density (g/cm3)  Exp. density (g/cm3)  Rel. density (%)  Mean grain size (mm)  Max fiber  length (μm)  TCTf  TCT  HCTf  HCT-1  HCT-2  HCT-3  1700, 9, 40  1750, 10, 30  1750, 14, 40  1750, 10, 30  1800, 10, 30  1900, 10, 30  12.25  13.42  10.89  12.33  12.33  12.33  11.87  13.42  10.32  10.29  10.79  10.72  96.7  100  94.8  83.5  87.5  86.9  2.27 1.4 2.37 1.2 0.67 0.4  -  0.77 0.4  -  98  -  124  -  -  -  \\x0c\", \"L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1369  Fig. 1. Low magnification SEM-images of  the chopped SiC Tyranno SA3 fibers (a) and of a fiber section (b).  method on the polished sections, at least 100 grains per specimen were measured. The apparent fiber length was measured on the polished sections perpendicular to the hot pressing direction, at least 100 fibers were analyzed. Vickers micro-hardness, HV1.0, was measured with a load of 9.81 N using a standard micro-hardness tester (Zwick 3212, Zwick, Ulm, Germany). Ten indentations were randomly placed onto the polished surfaces for each material; the diagonal of the indent was 30-33 μm, assuring the analysis of a representative volume of material. Young's modulus, E, was measured by the frequency method on 28 \\x02 8 \\x02 0.8 mm3 resonance specimens using a HP gain-phase analyzer. Fracture toughness, KIc, was evaluated using chevron-notched beams (CNB) in flexure. The test bars 25 \\x02 2 \\x02 2.5 mm3 (length by width by thickness) were notched with a 0.1 mm-thick diamond saw; the chevron-notched tip depth and average size length were about 0.12 mm and 0.80 mm of the bar thickness, respectively. The specimens were fractured using a semi-articulated silicon carbide in four-point fixture with a lower span of 20 mm and an upper span of 10 mm using screw-driven load frame (Instron mod. 6025, Instron, Illinois Tool Works Inc., Norwood, MA or Instron mod. 6025, High Wycombe, UK). The specimens were loaded with a crosshead speed of 0.05 mm/min. The “slice-model” equation of Munz et al. [30] was used to calculate KIc. On the same the flexural strength σ was machine and with the same fixture, 1200 1C on measured at room temperature and at test bars 25 \\x02 2 \\x02 2.5 mm3 (length by thickness by width), using a crosshead speed of 0.5 mm/min. The flexural strength at hightemperature was tested under flowing argon protective gas. Before the bending test, a soaking time of 18 min was set to reach thermal equilibrium. The thermal expansion coefficient CTE was measured up to 1300 1C under flowing argon with a 5 1C/min heating rate, using a dilatometer Netzsch mod. DIL E 402 (Netzsch, Geraetebau, Germany), on test bars 25 \\x02 2.5 \\x02 2 mm3 (length by thickness by width). The thermal conductivity up to 1900 1C was determined by measuring the thermal diffusivity with the laser-flash method and the specific heat under flowing argon. Thermogravimetrical analyses were carried out up to 1500 1C in air without forced gas flow with a heating rate of 10 1C/min, using the Simultaneous Thermal Analyzer STA 409 (Netzsch,  Geraetabau, Germany), on regular samples with dimensions 10-13 \\x02 2.5 \\x02 2 mm3. Before the test, all the specimens were cleaned in acetone and in ethanol and their dimensions measured. The weight of the specimens was measured before and after the exposure. The degradation after oxidation was evaluated by normalizing the weight gain, due to the oxide formation, on the surface area before exposure (%/cm2). initial Identical test bars were exposed for 5 min in a bottom loading furnace box when the maximum temperature of 1600 1C was reached. The microstructural modifications induced by oxidation were evaluated by x-ray diffraction and SEM-EDS analysis.  3. Results  3.1. Densification  Sintering cycles and density data relevant to TCTf and HCTf are summarized in Table 1, together with those of unreinforced TaCand HfC-materials. It can be noticed that for densification of fiber-reinforced composites, sintering temperatures were the same or lower than for the baseline matrices. The relative densities were expressed as the ratio between experimental and theoretical densities. The theoretical values nominal starting compositions and taking TaC density ¼ 13.90 g/ were calculated with the rule of the mixture, considering the cm3, HfC density ¼ 12.69 g/cm3, TaSi2 density ¼ 9.07 g/cm3 and SiC fibers density¼ 3.10 g/cm3. The contribution of low density secondary phases such as SiC or SiO2 was also considered for the calculation of the relative density. Their volumetric amount was estimated by image analysis on sintered samples. It is worthy to note that unreinforced hafnium carbide, HCT-1, HCT-2, and HCT-3, did not reach the full densification despite the sintering temperature being increased to 1750, 1800 and 1900 1C. The highest density, 10.79 g/cm3, was obtained in case of the sample sintered at 1800 1C (HCT-2).  3.2. Microstructure  3.2.1. TaC-10 TaSi2-15 SiC fibers (TCTf)  This composite was sintered at 1700 1C and reached a relative density of 97%. The x-ray diffraction pattern, carried out on the  \\x0c\", '1370  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  surface of the sample, in Fig. 2a shows only crystalline TaC (PDF # 65-0282) and TaSi2 (PDF # 38-0483) after sintering. The fractured and the polished sections are shown in Fig. 3; the matrix mean grain size is around 2 mm and no porosity was detected (Table 1). The fracture surface shows that fine TaC grains are intergranularly fractured whilst coarser TaC grains (4 2 mm) transgranularly fractured (Fig. 3a), in addition some fibers remain just half embedded in the matrix (as indicated by arrows) indicating weak bonds with the matrix. Fibers are  distributed almost isotropically in the section plane perpendicular to the hot pressing direction and their maximum length is around 180 mm, with a notable amount of fiber segments formed during milling (Fig. 3b). Large areas of TaSi2, recognizable as light gray phases with dimension 6-7 mm, were also observed in the matrix; occasionally these were intergranularly micro-cracked (inset in Fig. 3c). Other secondary phases scattered in the matrix were SiC/Si-C-O and SiO2 deriving both from the additive dissociation and oxygen contamination of the starting powder.  Fig. 2. X-ray diffraction pattern of TaC (a) and HfC (b) with fibers.  Fig.  3. TaC with  fibers  (a)  fracture  surface with  protruding  fibers  indicated  by  arrows,  (b)  fiber  distribution  and  detail  of  the  fiber/matrix  interface,  (c) microstructural  features of  the polished surface and detail of a TaSi2 micro-cracked phase, and (d)  typical shape of a fiber section.  \\x0c', 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1371  These phases were also found both at the fiber/matrix interface as visible in the inset of Fig. 3b. As for the reinforcement morphology, the fibers maintained the pristine appearance with a dark carbon-rich phase in the core and denser SiC grains in the rim (Fig. 3d).  3.2.2. HfC-10 TaSi2-15 SiC fibers (HCTf)  This composite was sintered at 1750 1C and reached a relative density of 95%. The x-ray diffraction pattern, carried out on the surface of the sample, in Fig. 2b shows only HfC as a crystalline phase. However HfC peaks do not exactly correspond to those of pure HfC (PDF # 39-1491), but are split and slightly shifted towards higher angles, owing to the formation of the (Hf,Ta)C solid solution, as found for similar composites [11,31]. Low intensity peaks visible at angles below 401 correspond to monoclinic hafnia. The fractured and polished sections in Fig. 4 show a fine 0.6 mm (Table microstructure, with grains around 1). The fracture surface evidences some protruding fibers, residual porosity, intergranular fracture of the squared HfC grains and transgranular fracture of TaSi2 which displays a ductile aspect (inset in Fig. 4a). As in the case of TCTf, the apparent maximum fiber length was below 200 μm (Fig. 4b). The matrix grains have a pure HfC core and a rim constituted by (Hf,Ta)C solution and HfSi2 has replaced TaSi2, as a consequence of cationic exchange between HfC grains and TaSi2 [11,31]. In addition, about 4 vol% of HfO2, with dimension up to 6-7 mm and white contrast, was recognized throughout the matrix and adjacent to the fibers, inset in Fig. 4b. The presence  of HfO2 was probably due to oxygen impurities in the starting HfC powder. At the fiber/matrix interface also for this composite, formation of SiC/SiCO and SiO2 dark pockets was noticed, inset in Fig. 4c. However, different from the previous case, most of the fiber body resulted notably changed upon densification, as it can be observed in Fig. 4c and d; the core was richer in carbon 1.5 μm thick, whilst the rim, about showed coarsened and dense SiC crystallites and no carbon pockets. The delimitation among the two zones was spotted by HfxSiy bright particles.  3.2.3. TaC and HfC matrices  The matrices (not shown) had fine microstructures with mean grain size o 2 μm. TaC-based matrix had fracture characteristics similar to TaC with fibers: no porosity, small grain intergranurarly fractured and big grain intragranularly fractured. Also the polished surface was similar presenting SiC/Si-C-O and SiO2 pockets and large areas of TaSi2. In the case of HfC-based matrix, higher amount of porosity, about 7%, SiC/SiCO, SiO2 dark pockets and HfO2 large zones were found.  3.3. Thermo-mechanical properties  The room temperature mechanical properties of TaCand HfC-composites are shown in Table 2. For comparison, available room temperature properties for unreinforced TaC and the most dense HfC matrix (HCT-2) are also reported. Vickers hardness among the TaC-based composites did not significantly change upon addition of SiC fibers and was  Fig. 4. HfC with fibers (a) fracture surface with arrows evidencing protruding fibers, (b) fiber distribution and detail of the fiber/matrix interface, (c) microstructural  features of  the polished surface, and (d)  typical shape of a fiber section.  \\x0c', \"1372  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  around 14 GPa. As for HfC-based composites, hardness is lower for unreinforced HfC (11 GPa), because of the higher content of residual porosity. Young's modulus of the composites containing fibers is slightly lower than typical values reported in the literature, around 490 GPa for both TaC and HfC, owing to soft silica-based phases and to little residual porosity, especially for HCTf. The fracture toughness of TaC-based composites resulted unchanged when fibers were added, while a slight improvement þ 30%. HfC was obtained for the HfC-based composite, toughness values are usually lower than those of TaC composites, because of the more brittle nature of the matrix [31]. As for the strength, a slight decrease of the average value is observed for TaC with the addition of fibers. On the contrary, among HfC-based composites, the flexural strength was higher for reinforced composites than for the baseline matrices. The high temperature flexural strength of the reinforced carbides was reduced compared to the room temperature value: \\x00 26% for TCTf and \\x00 23% for HCTf, similar to analogous carbides [31]. The linear CTE was similar for both ceramics containing fibers, 6.5-7 \\x02 10 \\x00 6 K \\x00 1, and not notably different from other transition metal carbides [32]. For the matrices, the CTE was notably higher only for HCT, again due to its residual porosity. The thermal conductivity increased for both materials up to 1000 1C and remained stable up to 1900 1C. Thermal conductivity of TaC-based composite was significantly higher than that of HfC-based one, especially above 500 1C (Table 3).  3.4. Oxidation tests  Thermogravimetrical analyses were carried out on TaC and HfC composites up to 1500 1C in air (Fig. 5). For HfC-based composites, the normalized weight gain per unit surface (ΔW/S) was similar for both reinforced and unreinforced materials, around 2.3%/cm2. TCT was strongly oxidized during the TG test with ΔW/S reaching 6.4%/cm2 and extensive oxide spalling was also observed, Fig. 5c. On the other hand, the oxidation behavior of TCTf was notably improved (with weight gain similar to that of HCT and HCTf, 2.6%/cm2), indicating that the oxidation of SiC fibers to silica exerted a beneficial effect in hindering oxygen penetration through the bulk and limiting spalling phenomena. According to x-ray diffraction, no difference in crystalline phases was detected in composites with or without fibers, i.e. Ta2O5 for TaC-based materials and HfO2 and Hf6Ta2O17 for HfC-based ones. The oxidized surfaces are shown in Fig. 6. The formed oxides showed the tendency to  become more compact with the addition of fibers (due to silica formation) compared to their respective matrices. On composites containing fibers, oxidation at 1600 1C for 5 min was performed too and SEM images of surface and cross section are reported in Fig. 7. The external appearance of the two samples is depicted in the insets of Fig. 7a and b, and it can be immediately seen that the TaC-based one resulted much more oxidized than the HfC-based one, but intact as compared to the TG test (compare Fig. 5). The external surface of TCTf is depicted in Fig. 7a and it is notably different from the surface of the ceramic after the thermogravimentric test in Fig. 6a, although the same crystalline phase was found. In this case, most of the fibers were consumed leaving silica-based pools and holes and melting of tantalum oxide occurred, in line with the Ta2O5-SiO2 eutectic at around 1560 1C [33]. On the other side, the external surfaces of HCTf after TG-test at 1500 1C (Fig. 6b) and elevator test at 1600 1C (Fig. 7b) were quite similar, with rough HfO2 and Hf6Ta2O17 grains and SiO2 in place of the fibers.  Table 3 Thermal conductivity (KTH) at 20, 500, 1000, 1500 and 1900 1C.  Label  KTH (W/m K)  20 1C  500 1C  1000 1C  1500 1C  1900 1C  TCTf  HCTf  27.8  20.7  33.6  24.7  37.1  26.0  37.2  26.0  37.2  26.5  Fig. 5. TGA plots of TaCand HfC-based ceramics during test at 1500 1C in 10 1C/min)  (heating  samples  after  pictures  of  the  rate  of  the  test:  air  and  (a) TCTf,  (b) HCTf,  (c) TCT, and (d) HCT.  Table 2 Thermo-mechanical properties: Vickers hardness (HV 1.0), elastic modulus (E), CNB fracture toughness (KIc), 4-pt bending strength at room temperature (σRT) and at 1200 1C (σ1200), and thermal expansion coefficient  (λ25-1300).  Label  TCTf  TCT  HCTf  HCT-2  HV 1.0 (GPa)  14.6 7 0.83 14.1 7 0.51 15.1 7 1.3 11.2 7 1.4  E (GPa)  4707 5  -  3737 5  -  (MPa \\x01 m1/2)  KIc  4.50 7 0.08 4.91 7 0.47 3.82 7 0.24 2.91 7 0.11  σRT (MPa)  4327 109 5067 145 3757 33 3237 48  σ1200 (MPa)  3207 9  -  2877 12  -  λ25-1300 (10  \\x00 6 K  \\x00 1)  6.66  7.63  7.01  9.99  \\x0c\", 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1373  Fig. 6. TCTf  (a) and HCTf  (b) surfaces after  the TGA test.  The cross section of the oxidized specimens (Fig. 7c and d), evidenced that about 290 μm of material underwent modification in TCTf, whilst only 145 μm in HCTf. In TCTf, Fig. 7c, the first 22 μm were composed by coarse Ta2O5 grains with discontinuous zones of SiO2 where tantalum oxide was dispersed in, then a 270 μm of fine grained tantalum oxide layer with holes left by the fiber oxidation and veined by silica glass followed. In the bulk, traces of oxygen were detected in the matrix, whilst the fiber was pure SiC. In HCTf, Fig. 7d, the outermost 80 μm were composed by fine HfO2 and rounded Hf6Ta2O17 grains with partially oxidized SiC fibers from where discontinuous SiO2 glass spilled out. In this layer large cavities formed owing to volatile products formation, as previously outlined for HfCand TaSi2-containing composites [34]. The following 65 μm were composed by dense HfO2 zones dispersed in porous HfO2 with similar grain size. It is probable that the dense areas correspond to the former hafnia regions already present in the sintered microstructure (Fig. 4b), whilst the porous region are just the oxidation result of the carbide. In this scale the fibers were already quite compromised by the oxidation, but carbon peak was still detected by EDS. Moving further down, HfC contoured by Hf-O-C constituted the matrix and the fiber showed preferential oxidation of the Hf-Si phase which originally separated the core and rim regions (compare Fig. 4b).  4. Discussion: effect of SiC fiber addition on the properties of carbides  The main purpose of fiber addition is to get an improvement in fracture toughness for brittle UHTC matrices. This goal was reasonably achieved in the case of previously studied ZrB2 [35], but not for these carbides. The reason for the poor efficacy of fibers in these carbide matrices is still under investigation. In order to get an insight into the fracture mode of these composites, a 98.1 N indentation was introduced onto the polished surfaces, Fig. 8. Most of the times, both in the case of TaC and in the case of HfC, the fibers were cut across, but in many cases the crack could also proceed along the fiber/matrix boundary, if this was arranged along the longitudinal direction (insets in Fig. 8). On the other hand, looking at the fracture surfaces after bending strength tests, it can be seen that during fracture propagation, some fibers  K composite I c  ¼ K matr ix I c  þ ΔK pinning I c  I c  ¼ K matr ix I c  þ ΔK res:stress I c  were pulled apart. For previously studied reinforced composites, crack pinning was supposed to be the dominant mechanism justifying the observed increases in toughness [35] because crack deflection and crack bridging mechanisms were negligible, due to very strong interface. Moreover, residual stress mechanisms were supposed to give a negative contribution, due to mismatch between elastic moduli and CTE of the two phases. Since for TaC composite K composite we conclude that if any positive toughening effect is exerted by the fibers, it is counterbalanced by negative contributions from residual stresses, such that the total effect is null. Preliminary calculations summing each specific contribution to toughness, i.e. , did not result well in agreement with the experimental results, with the theoretical values always exceeding the experimental ones. This could be due to several reasons, for example the introduction of the input values required by the models (elastic constant, coefficient of thermal expansion and fiber strength) [36,37] could not be coherent with the actual values of the composite after sintering. Increase of residual stresses or, most probably, decrease of tensile fiber strength due to chemical alteration upon thermal treatments could occur. This aspect deserves surely further investigation and, in particular, the actual properties of the fibers in the ceramic should be measured. In the case of HfC a modest increase was recorded, however, this was attributed more to the increase in the final density of the matrix, which in turn determined an increase of its toughness, rather than effective toughening mechanisms activated by the fibers. This hypothesis is corroborated by previous results obtained on dense HfC materials, for which the fracture toughness without any reinforcement was 3.6 MPa m1/2 [10]. Flexural strength is another property affected by the fiber addition. It should be noticed that for TaC-based matrix, room temperature strength data are quite scattered. Critical defects can be large TaSi2 pockets, especially when micro-cracked, as observed in Fig. 3c. Moreover, being TaC a very hard and brittle ceramic, surface defects due to machining may also affect the strength. A notable data scattering for TaC was reported also by Zhang et al. [9]. With the addition of fibers, the strength of this hard carbide did not statistically change. A slight decrease of the average values strength was observed, accompanied by a decrease of the standard deviation. Very  \\x0c', '1374  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  Fig. 7. External surface (top) and cross section (bottom) of (a, c) TCTf and (b, d) HCTf after the oxidation test at 1600 1C for 5 min in air. The insets in (a and b)  show the appearance of  the specimens upon oxidation.  likely, the dimension of the critical defects generated by fibers was not significantly different from pristine ones. On the contrary, for HfC, matrix strength values are rather low, due to incomplete sintering. Indeed, typical values measured on similar unreinforced matrices range from 460 to 540 MPa for hot pressed or pressureless sintered matrices with TaSi2 or MoSi2 as sintering agents [14,20,31]. With the addition of fibers, a slight improvement was detected which is due to the increase of final density. In this case, SiC fibers acted more as a sintering aid rather than a strengthening phase. This is clearly  demonstrated by microstructural analyses; indeed large porous regions with hafnia particles observed in the matrix (HCT-2) were replaced by dense mixed SiC-HfO2 regions which were just localized in the proximity of SiC fibers, as shown in Fig. 2b and c. For the same reason, an increase of hardness was observed in HCTf compared to HCT-2 (Table 2). Other properties, such as thermal conductivity, are expected to be affected by the addition of SiC fibers. Although we did not measure the thermal conductivity of the respective matrices, unpublished KTH measured on TaC and HfC matrices with MoSi2  \\x0c', 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1375  as sintering aid showed the same trend as the composites object of this publication, i.e. thermal conductivity increase with the temperature and higher values for TaC-based compounds. For those analogous ceramics with MoSi2 the conductivity was similar to the reinforced composites up to 500 1C, but at 1900 1C the unreinforced composites had a conductivity up to 10 W/m K, higher than the SiC-fiber reinforced ones, since at high temperature, the thermal conductivity of HfC and TaC increases, whilst it decreases for SiC. Discussing the effects of SiC fiber on the overall thermal conductivity is not straightforward because we do not thermal conductivity up to 1900 1C, but we know the fiber can assume that the introduction of SiC fiber in TaC or HfC matrix causes a conductivity decrease at high temperature. Moreover, microstructural changes due to exposition at high temperature, such as SiC active oxidation in the subsurface layer resulting in the creation of voids could also affect this property. HCTf section after thermal conductivity test is shown in Fig. 9a and subsurface voids are pointed out in the inset. Fig. 9b and c show the surface covered by SiC due to the graphite-based painting used for the test. However we can assess that the lower values for HCTf are due both to the higher porosity of its matrix (about 7%) as compared to the dense TCTf and to the intrinsic lower conductivity of HfC.  As for the oxidation behavior, from SEM analyses reported in Fig. 7 it is clear that, both for TaCand HfC-composites, 15 vol% of fiber were not enough to form a continuous silica layer and therefore to form an efficient barrier against oxygen penetration at 1600 1C. There can be several reasons for such behavior: there is not enough silica to fill all the volume expansion associated with the formation of the oxides, or the thermodynamic involved causes volatilization of silica to gaseous SiO on the surface. Though, we can confirm that the introduction of SiC fiber improved the oxidation resistance of TaC, as evidenced in Fig. 5. In summary, the addition of SiC fibers did not produce any measurable benefit in terms of fracture toughness or strength. Compared to available literature data on composites HfC and TaC reinforced with SiC particles, it can be observed that TaC is reinforced more efficiently by SiC particles than SiC fibers. For instance, in the work of Liu et al. [38], addition of 10 vol% SiC particles resulted in a 3-pt bending strength of 660 MPa, and fracture toughness approaching 6 MPa m1/2. On the contrary for HfC, addition of SiC particles resulted in toughness values of 2.2-2.8 MPa m1/2 (i.e. even lower than upon addition of fibers) and similar 4-pt bending strength values (396 MPa) [39]. A significant contribution from SiC fiber mainly invested the densification behavior, especially in the case of HfC. If on one  Fig. 8. Crack path crossing the fibers generated by a 98.1 N indentation onto TaC (a) and HfC (b) with fibers and details of  the advancing crack along the  fiber/matrix boundary.  Fig. 9. HCTf after thermal conductivity test: (a) section and detail of the voids in the subsurface, (b) SiC grains on the surface, and (c) detail of the surface showing  HfC grains among SiC ones.  \\x0c', 'side the fiber released some silica-based phases that were beneficial for obtaining a denser matrix, on the other side a too strong interaction occurred between the fiber and matrix causing notable fibers modification of morphology and properties and not allowing efficient toughening mechanisms.  5. Conclusions  TaCand HfC-based composites containing 15 vol% short SiC Tyranno SA3 fibers were produced by hot pressing at 1700-1750 1C. TaSi2 was selected as sintering additive to lower the maximum sintering temperature and preserve SiC fibers. The addition of SiC fibers enabled densification of carbides at temperatures even lower than the respective matrices, but no benefits in terms of fracture toughness or strength were observed. The better densification of HfC generated a modest increase of its properties, whilst those of TaC were unaffected. The mechanical properties of the composites were mainly dictated by the matrix properties being higher for TaC than for HfC based ceramics. On the other hand, the oxidation resistance of TaC was somewhat improved with fiber addition, reaching the performance of more oxidation resistant HfCbased composites, either reinforced or unreinforced.  Acknowledgments  D. Dalle Fabbriche, C. Capiani are gratefully acknowledged for their technical support. Stefano Guicciardi is acknowledged for useful discussion on the mechanical properties.  References  [1] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, William  Andrew Publishing, Norwich, NY, 2001.  [2] L.E. Toth, Transition metal carbides and nitrides,  in: J.L. Margrave (Ed.),  Refractory Materials,  a Series  of Monographs, Academic Press, New  York, 1967, pp. 6-10.  [3] E.K. 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},{
  "_id": 130,
  "PDF": "Microstructure,mechanica lproperties and oxidation behavior of TaC-and HfC-based materials containing short SiC fiber.pdf",
  "Text": "[\"Available online at www.sciencedirect.com  Ceramics International 41 (2015) 1367-1377  CERAMICS  INTERNATIONAL  www.elsevier.com/locate/ceramint  Microstructure, mechanical properties and oxidation behavior of TaCand HfC-based materials containing short SiC fiber  L. Pienti, L. Silvestronin, E. Landi, C. Melandri, D. Sciti  CNR-ISTEC,  Institute of Science and Technology for Ceramics, Via Granarolo 64,  I-48018 Faenza,  Italy  Received 5 June 2014;  received in revised form 1 September 2014; accepted 12 September 2014  Available online 19 September 2014  Abstract  TaCand HfC-based composites containing 15 vol% short SiC fibers as reinforcing phase were produced by hot pressing at 1700-1750 1C. Suitable sintering additives were selected in order to get a full densification at temperatures well-tolerated by SiC fibers. Microstructural  characterization outlined a strong interaction between the carbide matrix and SiC fibers. The following properties were evaluated: Vickers hardness, fracture toughness by chevron notched beams, elastic modulus, 4-pt bending strength at room temperature and at 1200 1C, linear CTE up to 1300 1C in Ar, and thermal conductivity up to 1900 1C in Ar. Oxidation tests in air were carried out with thermogravimetric analysis up to 1500 1C and in a bottom up furnace at 1600 1C. TaC-based materials generally possessed higher thermo-mechanical properties compared to HfC, but lower oxidation resistance. Properties of fiber-reinforced composites were compared to those of unreinforced ones.  & 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  Keywords: Fibers; TaC; HfC; Thermo-mechanical properties  1.  Introduction  Among advanced ceramics, tantalumand hafniumbased refractory carbides are promising materials for applications in the aerospace and energy sectors. These materials combine the physical properties of ceramic and the electronic properties of metals. They are stiff, with Young's modulus values competing with those of SiC, and have good thermal conductivity, permitting heat to be drawn away from the superheated surfaces. Furthermore, they have high melting points (4 3900 1C) and good chemical resistance. These properties make TaC and HfC interesting systems for extremely hightemperature applications such as rocket nozzles, scramjet components, and future generation thrusters that require resistance to severe combustion conditions [1-3]. Mono l i th ic TaC and HfC are hard to dens ify due to the ir s trong cova len t bonds and low se lf-d if fus ion coeffic ien ts [1-3] . Ho t press ing , spark p lasma s in ter ing and pressureless s in ter ing have been of ten assoc ia ted to the use o f  nCorresponding author. Fax: þ 39 0546 46381. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  http://dx.doi.org/10.1016/j.ceramint.2014.09.070  0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  add i t ives to imp rove dens ifica t ion and m i t iga te the severe s in ter ing tempe ra tures [4-21] . Sma l l amoun ts o f me ta ls such as Fe , Mn , Co , and N i , lower the tempera tu re needed for dens ifica t ion [5-7] ; a l terna t ive s in ter ing add i t ives such as B , C , and B4C , remove surface ox ides on the s ta r t ing 1900-2300 1C par t ic les enab l ing dens ifica t ion a t [8 ,9 ] . Me ta l d is i lic ides , such as TaS i2 and MoS i2 have shown to be effec t ive add i t ives a l low ing dens ifica t ion of TaC and HfC a t tempera tures be tween 1750 and 1950 1C [11] e i ther by ho t press ing or pressure less s in ter ing [21] . Reduc ing par t ic le s ize of the powders has a lso been shown to dens ifica t ion . TaC and HfC þ 5 vo l% MoS i2 improve the w ith re la t ive dens i t ies above 90% have been ob ta ined s tar t ing from syn thes ized u ltrafine powders [10 ] . Recen t ly , spark p lasma s in ter ing was a lso used to conso l ida te Ta0 .8H f0 .2C us ing MoS i2 and TaS i2 as s in ter ing a ids a t 1650 1C under 30 MPa [19] . Published values for the thermo-mechanical properties of these carbides are still scarce. Hardness of 21-27 GPa and fracture toughness of 3.5 MPa m1/2 were measured by Balani et al. on vacuum plasma sprayed TaC [22]. TaC Young's modulus value of 550 GPa was measured by Lopez-de-la-Torre et al. [23]. For hot  \\x0c\", \"1368  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  pressed HfC-based composites with 15 vol% MoSi2, hardness up to 19 GPa was obtained [14], and mechanical strength of 450- 465 MPa and Young's modulus of 415-434 GPa for pressureless composites with 5-10 vol% of the same sintering additive were measured [16]. There are no recent reports on thermal properties like linear CTE up to 1300 1C in Ar, thermal conductivity up to 1900 1C in Ar. One drawback of tantalum and hafnium carbides is the poor oxidation resistance [1-3]. TaC oxidizes to Ta2O5 which has a relatively low melting point, 1800 1C [24]. According to previous studies on the high temperature oxidation of HfC-TaC compo[25,26], at To 1800 1C the outermost oxide layer sites formed was constituted of Ta2O5, while at T4 1800 1C it was made of a mixture of Ta2Hf6O19 and HfO2. HfC oxidizes to monoclinic HfO2, one of the most stable refractory compounds with a melting point of 2800 1C [24]. Oxidation tests carried out between 1400 and 2060 1C [27] and arc jet tests between 2400 and 2700 1C [28] confirmed the excellent oxidation response in severe condition of monolithic HfC. Similarly, HfC composites with MoSi2 at 5 vol% showed a good thermal stability when tested in an arc jet facility up to 2400 1C [29]. In this work the effect of SiC fibers on microstructure, thermomechanical properties and oxidation resistance of TaCand HfCbased composites was analyzed. TaSi2 has been selected as a sintering additive for the hot pressing process as it allows densification at relatively low temperature, compatible with the stability limit of SiC fibers. Interaction between the carbide matrix and SiC fibers was studied. The following properties were measured: Vickers hardness, fracture toughness by chevron notched beams, elastic modulus, 4-pt bending strength at room temperature and at 1200 1C, linear CTE up to 1300 1C in Ar and thermal conductivity up to 1900 1C in Ar. The oxidation behavior was tested through thermogravimetric balance up to 1500 1C and in a bottom up loading furnace at 1600 1C.  2. Experimental procedure  produced  starting  from commercial  The composites were raw materials: \\x0f TaC (Treibacher Industrie AG, Althofen, Austria), mean particle size 1.1 μm; \\x0f HfC (Cerac Inc., Milwaukee, WI, USA), particle size range \\x0f TaSi2 (ABCR GmbH & Co., Karlsruhe, Germany), particle 0.2-1.5 mm; size o 45 mm;  \\x0f SiC fibers (Tyranno SA3, UBE Industries, Ltd.) diameter 7.5-10 mm, composition wt% Si:C ¼ 67:31, oxygen o 1 wt%, and aluminum o 2 wt%.  labeled as reported in Table 1,  The following compositions, were prepared (vol%): \\x0f TaC þ 10 TaSi2 þ 15 chopped SiC fibers (TCTf); \\x0f HfC þ 10 TaSi2 þ 15 chopped SiC fibers (HCTf).  For the sake of comparison, the respective matrices were e.g.: TaC þ 10 vol% TaSi2 also produced and characterized, (TCT) and HfC þ 10 vol% TaSi2 (HCT). The powders were dispersed in absolute ethanol by simultaneously applying magnetic stirring and ultrasonication. For the reinforced composites, fibers were in house chopped. Low magnification SEM-images of the chopped SiC fibers and of a fiber section are reported in Fig. 1. The powder mixtures were further ball milled in absolute ethanol for 24 h, at 130 rpm, with SiC milling media. A fixed powder/milling media/solvent weight amount (1:1:1) was adopted. After milling, the slurries were dried in a rotary evaporator, sieved and shaped in four centimeters diameter pellets. The pellets with fibers were debonded at 500 1C for 1 h and all the samples were hot 1700-1900 1C range pressed at temperatures in the in low vacuum (100 Pa) using an induction-heated graphite die with a uniaxial pressure of 30 MPa, that was increased to 40 MPa during the isothermal stage; free cooling followed. For each composition, the maximum sintering temperature was set on the basis of the shrinkage curve, Table 1. After sintering, the bulk densities were measured by Archimedes' method. Crystalline phases were identified by x-ray diffraction (Bruker D8 Advance, Bruker, Karlsruhe, Germany). The microstructure was analyzed using a scanning electron microscope (FE-SEM, Carl Zeiss Sigma NTS GmbH, Oberkochen, Germany) and an energy-dispersive spectroscopy (EDS, X-Act, INCA Energy 300, Oxford Instruments, Abingdon, UK) on fractured and polished surfaces. Sections of the sintered materials were cut perpendicularly to the hot pressing direction and polished with diamond paste to 0.25 mm. Mean matrix grain size, amount of porosity and amount of secondary phases were determined on micrographs of polished sections using image analysis (Image-Pro Analyzer 7.0, Media Cybernetics, Silver Spring, MD, Rockville, USA). The mean grain size of the matrix grains was calculated by the circle  Table 1  List of  the samples:  label, sintering conditions (TMAX, dwell  time, pressure), density, matrix mean grain size and maximum fiber  length.  Label  Sintering (1C, min, MPa)  Th. density (g/cm3)  Exp. density (g/cm3)  Rel. density (%)  Mean grain size (mm)  Max fiber  length (μm)  TCTf  TCT  HCTf  HCT-1  HCT-2  HCT-3  1700, 9, 40  1750, 10, 30  1750, 14, 40  1750, 10, 30  1800, 10, 30  1900, 10, 30  12.25  13.42  10.89  12.33  12.33  12.33  11.87  13.42  10.32  10.29  10.79  10.72  96.7  100  94.8  83.5  87.5  86.9  2.27 1.4 2.37 1.2 0.67 0.4  -  0.77 0.4  -  98  -  124  -  -  -  \\x0c\", \"L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1369  Fig. 1. Low magnification SEM-images of  the chopped SiC Tyranno SA3 fibers (a) and of a fiber section (b).  method on the polished sections, at least 100 grains per specimen were measured. The apparent fiber length was measured on the polished sections perpendicular to the hot pressing direction, at least 100 fibers were analyzed. Vickers micro-hardness, HV1.0, was measured with a load of 9.81 N using a standard micro-hardness tester (Zwick 3212, Zwick, Ulm, Germany). Ten indentations were randomly placed onto the polished surfaces for each material; the diagonal of the indent was 30-33 μm, assuring the analysis of a representative volume of material. Young's modulus, E, was measured by the frequency method on 28 \\x02 8 \\x02 0.8 mm3 resonance specimens using a HP gain-phase analyzer. Fracture toughness, KIc, was evaluated using chevron-notched beams (CNB) in flexure. The test bars 25 \\x02 2 \\x02 2.5 mm3 (length by width by thickness) were notched with a 0.1 mm-thick diamond saw; the chevron-notched tip depth and average size length were about 0.12 mm and 0.80 mm of the bar thickness, respectively. The specimens were fractured using a semi-articulated silicon carbide in four-point fixture with a lower span of 20 mm and an upper span of 10 mm using screw-driven load frame (Instron mod. 6025, Instron, Illinois Tool Works Inc., Norwood, MA or Instron mod. 6025, High Wycombe, UK). The specimens were loaded with a crosshead speed of 0.05 mm/min. The “slice-model” equation of Munz et al. [30] was used to calculate KIc. On the same the flexural strength σ was machine and with the same fixture, 1200 1C on measured at room temperature and at test bars 25 \\x02 2 \\x02 2.5 mm3 (length by thickness by width), using a crosshead speed of 0.5 mm/min. The flexural strength at hightemperature was tested under flowing argon protective gas. Before the bending test, a soaking time of 18 min was set to reach thermal equilibrium. The thermal expansion coefficient CTE was measured up to 1300 1C under flowing argon with a 5 1C/min heating rate, using a dilatometer Netzsch mod. DIL E 402 (Netzsch, Geraetebau, Germany), on test bars 25 \\x02 2.5 \\x02 2 mm3 (length by thickness by width). The thermal conductivity up to 1900 1C was determined by measuring the thermal diffusivity with the laser-flash method and the specific heat under flowing argon. Thermogravimetrical analyses were carried out up to 1500 1C in air without forced gas flow with a heating rate of 10 1C/min, using the Simultaneous Thermal Analyzer STA 409 (Netzsch,  Geraetabau, Germany), on regular samples with dimensions 10-13 \\x02 2.5 \\x02 2 mm3. Before the test, all the specimens were cleaned in acetone and in ethanol and their dimensions measured. The weight of the specimens was measured before and after the exposure. The degradation after oxidation was evaluated by normalizing the weight gain, due to the oxide formation, on the surface area before exposure (%/cm2). initial Identical test bars were exposed for 5 min in a bottom loading furnace box when the maximum temperature of 1600 1C was reached. The microstructural modifications induced by oxidation were evaluated by x-ray diffraction and SEM-EDS analysis.  3. Results  3.1. Densification  Sintering cycles and density data relevant to TCTf and HCTf are summarized in Table 1, together with those of unreinforced TaCand HfC-materials. It can be noticed that for densification of fiber-reinforced composites, sintering temperatures were the same or lower than for the baseline matrices. The relative densities were expressed as the ratio between experimental and theoretical densities. The theoretical values nominal starting compositions and taking TaC density ¼ 13.90 g/ were calculated with the rule of the mixture, considering the cm3, HfC density ¼ 12.69 g/cm3, TaSi2 density ¼ 9.07 g/cm3 and SiC fibers density¼ 3.10 g/cm3. The contribution of low density secondary phases such as SiC or SiO2 was also considered for the calculation of the relative density. Their volumetric amount was estimated by image analysis on sintered samples. It is worthy to note that unreinforced hafnium carbide, HCT-1, HCT-2, and HCT-3, did not reach the full densification despite the sintering temperature being increased to 1750, 1800 and 1900 1C. The highest density, 10.79 g/cm3, was obtained in case of the sample sintered at 1800 1C (HCT-2).  3.2. Microstructure  3.2.1. TaC-10 TaSi2-15 SiC fibers (TCTf)  This composite was sintered at 1700 1C and reached a relative density of 97%. The x-ray diffraction pattern, carried out on the  \\x0c\", '1370  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  surface of the sample, in Fig. 2a shows only crystalline TaC (PDF # 65-0282) and TaSi2 (PDF # 38-0483) after sintering. The fractured and the polished sections are shown in Fig. 3; the matrix mean grain size is around 2 mm and no porosity was detected (Table 1). The fracture surface shows that fine TaC grains are intergranularly fractured whilst coarser TaC grains (4 2 mm) transgranularly fractured (Fig. 3a), in addition some fibers remain just half embedded in the matrix (as indicated by arrows) indicating weak bonds with the matrix. Fibers are  distributed almost isotropically in the section plane perpendicular to the hot pressing direction and their maximum length is around 180 mm, with a notable amount of fiber segments formed during milling (Fig. 3b). Large areas of TaSi2, recognizable as light gray phases with dimension 6-7 mm, were also observed in the matrix; occasionally these were intergranularly micro-cracked (inset in Fig. 3c). Other secondary phases scattered in the matrix were SiC/Si-C-O and SiO2 deriving both from the additive dissociation and oxygen contamination of the starting powder.  Fig. 2. X-ray diffraction pattern of TaC (a) and HfC (b) with fibers.  Fig.  3. TaC with  fibers  (a)  fracture  surface with  protruding  fibers  indicated  by  arrows,  (b)  fiber  distribution  and  detail  of  the  fiber/matrix  interface,  (c) microstructural  features of  the polished surface and detail of a TaSi2 micro-cracked phase, and (d)  typical shape of a fiber section.  \\x0c', 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1371  These phases were also found both at the fiber/matrix interface as visible in the inset of Fig. 3b. As for the reinforcement morphology, the fibers maintained the pristine appearance with a dark carbon-rich phase in the core and denser SiC grains in the rim (Fig. 3d).  3.2.2. HfC-10 TaSi2-15 SiC fibers (HCTf)  This composite was sintered at 1750 1C and reached a relative density of 95%. The x-ray diffraction pattern, carried out on the surface of the sample, in Fig. 2b shows only HfC as a crystalline phase. However HfC peaks do not exactly correspond to those of pure HfC (PDF # 39-1491), but are split and slightly shifted towards higher angles, owing to the formation of the (Hf,Ta)C solid solution, as found for similar composites [11,31]. Low intensity peaks visible at angles below 401 correspond to monoclinic hafnia. The fractured and polished sections in Fig. 4 show a fine 0.6 mm (Table microstructure, with grains around 1). The fracture surface evidences some protruding fibers, residual porosity, intergranular fracture of the squared HfC grains and transgranular fracture of TaSi2 which displays a ductile aspect (inset in Fig. 4a). As in the case of TCTf, the apparent maximum fiber length was below 200 μm (Fig. 4b). The matrix grains have a pure HfC core and a rim constituted by (Hf,Ta)C solution and HfSi2 has replaced TaSi2, as a consequence of cationic exchange between HfC grains and TaSi2 [11,31]. In addition, about 4 vol% of HfO2, with dimension up to 6-7 mm and white contrast, was recognized throughout the matrix and adjacent to the fibers, inset in Fig. 4b. The presence  of HfO2 was probably due to oxygen impurities in the starting HfC powder. At the fiber/matrix interface also for this composite, formation of SiC/SiCO and SiO2 dark pockets was noticed, inset in Fig. 4c. However, different from the previous case, most of the fiber body resulted notably changed upon densification, as it can be observed in Fig. 4c and d; the core was richer in carbon 1.5 μm thick, whilst the rim, about showed coarsened and dense SiC crystallites and no carbon pockets. The delimitation among the two zones was spotted by HfxSiy bright particles.  3.2.3. TaC and HfC matrices  The matrices (not shown) had fine microstructures with mean grain size o 2 μm. TaC-based matrix had fracture characteristics similar to TaC with fibers: no porosity, small grain intergranurarly fractured and big grain intragranularly fractured. Also the polished surface was similar presenting SiC/Si-C-O and SiO2 pockets and large areas of TaSi2. In the case of HfC-based matrix, higher amount of porosity, about 7%, SiC/SiCO, SiO2 dark pockets and HfO2 large zones were found.  3.3. Thermo-mechanical properties  The room temperature mechanical properties of TaCand HfC-composites are shown in Table 2. For comparison, available room temperature properties for unreinforced TaC and the most dense HfC matrix (HCT-2) are also reported. Vickers hardness among the TaC-based composites did not significantly change upon addition of SiC fibers and was  Fig. 4. HfC with fibers (a) fracture surface with arrows evidencing protruding fibers, (b) fiber distribution and detail of the fiber/matrix interface, (c) microstructural  features of  the polished surface, and (d)  typical shape of a fiber section.  \\x0c', \"1372  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  around 14 GPa. As for HfC-based composites, hardness is lower for unreinforced HfC (11 GPa), because of the higher content of residual porosity. Young's modulus of the composites containing fibers is slightly lower than typical values reported in the literature, around 490 GPa for both TaC and HfC, owing to soft silica-based phases and to little residual porosity, especially for HCTf. The fracture toughness of TaC-based composites resulted unchanged when fibers were added, while a slight improvement þ 30%. HfC was obtained for the HfC-based composite, toughness values are usually lower than those of TaC composites, because of the more brittle nature of the matrix [31]. As for the strength, a slight decrease of the average value is observed for TaC with the addition of fibers. On the contrary, among HfC-based composites, the flexural strength was higher for reinforced composites than for the baseline matrices. The high temperature flexural strength of the reinforced carbides was reduced compared to the room temperature value: \\x00 26% for TCTf and \\x00 23% for HCTf, similar to analogous carbides [31]. The linear CTE was similar for both ceramics containing fibers, 6.5-7 \\x02 10 \\x00 6 K \\x00 1, and not notably different from other transition metal carbides [32]. For the matrices, the CTE was notably higher only for HCT, again due to its residual porosity. The thermal conductivity increased for both materials up to 1000 1C and remained stable up to 1900 1C. Thermal conductivity of TaC-based composite was significantly higher than that of HfC-based one, especially above 500 1C (Table 3).  3.4. Oxidation tests  Thermogravimetrical analyses were carried out on TaC and HfC composites up to 1500 1C in air (Fig. 5). For HfC-based composites, the normalized weight gain per unit surface (ΔW/S) was similar for both reinforced and unreinforced materials, around 2.3%/cm2. TCT was strongly oxidized during the TG test with ΔW/S reaching 6.4%/cm2 and extensive oxide spalling was also observed, Fig. 5c. On the other hand, the oxidation behavior of TCTf was notably improved (with weight gain similar to that of HCT and HCTf, 2.6%/cm2), indicating that the oxidation of SiC fibers to silica exerted a beneficial effect in hindering oxygen penetration through the bulk and limiting spalling phenomena. According to x-ray diffraction, no difference in crystalline phases was detected in composites with or without fibers, i.e. Ta2O5 for TaC-based materials and HfO2 and Hf6Ta2O17 for HfC-based ones. The oxidized surfaces are shown in Fig. 6. The formed oxides showed the tendency to  become more compact with the addition of fibers (due to silica formation) compared to their respective matrices. On composites containing fibers, oxidation at 1600 1C for 5 min was performed too and SEM images of surface and cross section are reported in Fig. 7. The external appearance of the two samples is depicted in the insets of Fig. 7a and b, and it can be immediately seen that the TaC-based one resulted much more oxidized than the HfC-based one, but intact as compared to the TG test (compare Fig. 5). The external surface of TCTf is depicted in Fig. 7a and it is notably different from the surface of the ceramic after the thermogravimentric test in Fig. 6a, although the same crystalline phase was found. In this case, most of the fibers were consumed leaving silica-based pools and holes and melting of tantalum oxide occurred, in line with the Ta2O5-SiO2 eutectic at around 1560 1C [33]. On the other side, the external surfaces of HCTf after TG-test at 1500 1C (Fig. 6b) and elevator test at 1600 1C (Fig. 7b) were quite similar, with rough HfO2 and Hf6Ta2O17 grains and SiO2 in place of the fibers.  Table 3 Thermal conductivity (KTH) at 20, 500, 1000, 1500 and 1900 1C.  Label  KTH (W/m K)  20 1C  500 1C  1000 1C  1500 1C  1900 1C  TCTf  HCTf  27.8  20.7  33.6  24.7  37.1  26.0  37.2  26.0  37.2  26.5  Fig. 5. TGA plots of TaCand HfC-based ceramics during test at 1500 1C in 10 1C/min)  (heating  samples  after  pictures  of  the  rate  of  the  test:  air  and  (a) TCTf,  (b) HCTf,  (c) TCT, and (d) HCT.  Table 2 Thermo-mechanical properties: Vickers hardness (HV 1.0), elastic modulus (E), CNB fracture toughness (KIc), 4-pt bending strength at room temperature (σRT) and at 1200 1C (σ1200), and thermal expansion coefficient  (λ25-1300).  Label  TCTf  TCT  HCTf  HCT-2  HV 1.0 (GPa)  14.6 7 0.83 14.1 7 0.51 15.1 7 1.3 11.2 7 1.4  E (GPa)  4707 5  -  3737 5  -  (MPa \\x01 m1/2)  KIc  4.50 7 0.08 4.91 7 0.47 3.82 7 0.24 2.91 7 0.11  σRT (MPa)  4327 109 5067 145 3757 33 3237 48  σ1200 (MPa)  3207 9  -  2877 12  -  λ25-1300 (10  \\x00 6 K  \\x00 1)  6.66  7.63  7.01  9.99  \\x0c\", 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1373  Fig. 6. TCTf  (a) and HCTf  (b) surfaces after  the TGA test.  The cross section of the oxidized specimens (Fig. 7c and d), evidenced that about 290 μm of material underwent modification in TCTf, whilst only 145 μm in HCTf. In TCTf, Fig. 7c, the first 22 μm were composed by coarse Ta2O5 grains with discontinuous zones of SiO2 where tantalum oxide was dispersed in, then a 270 μm of fine grained tantalum oxide layer with holes left by the fiber oxidation and veined by silica glass followed. In the bulk, traces of oxygen were detected in the matrix, whilst the fiber was pure SiC. In HCTf, Fig. 7d, the outermost 80 μm were composed by fine HfO2 and rounded Hf6Ta2O17 grains with partially oxidized SiC fibers from where discontinuous SiO2 glass spilled out. In this layer large cavities formed owing to volatile products formation, as previously outlined for HfCand TaSi2-containing composites [34]. The following 65 μm were composed by dense HfO2 zones dispersed in porous HfO2 with similar grain size. It is probable that the dense areas correspond to the former hafnia regions already present in the sintered microstructure (Fig. 4b), whilst the porous region are just the oxidation result of the carbide. In this scale the fibers were already quite compromised by the oxidation, but carbon peak was still detected by EDS. Moving further down, HfC contoured by Hf-O-C constituted the matrix and the fiber showed preferential oxidation of the Hf-Si phase which originally separated the core and rim regions (compare Fig. 4b).  4. Discussion: effect of SiC fiber addition on the properties of carbides  The main purpose of fiber addition is to get an improvement in fracture toughness for brittle UHTC matrices. This goal was reasonably achieved in the case of previously studied ZrB2 [35], but not for these carbides. The reason for the poor efficacy of fibers in these carbide matrices is still under investigation. In order to get an insight into the fracture mode of these composites, a 98.1 N indentation was introduced onto the polished surfaces, Fig. 8. Most of the times, both in the case of TaC and in the case of HfC, the fibers were cut across, but in many cases the crack could also proceed along the fiber/matrix boundary, if this was arranged along the longitudinal direction (insets in Fig. 8). On the other hand, looking at the fracture surfaces after bending strength tests, it can be seen that during fracture propagation, some fibers  K composite I c  ¼ K matr ix I c  þ ΔK pinning I c  I c  ¼ K matr ix I c  þ ΔK res:stress I c  were pulled apart. For previously studied reinforced composites, crack pinning was supposed to be the dominant mechanism justifying the observed increases in toughness [35] because crack deflection and crack bridging mechanisms were negligible, due to very strong interface. Moreover, residual stress mechanisms were supposed to give a negative contribution, due to mismatch between elastic moduli and CTE of the two phases. Since for TaC composite K composite we conclude that if any positive toughening effect is exerted by the fibers, it is counterbalanced by negative contributions from residual stresses, such that the total effect is null. Preliminary calculations summing each specific contribution to toughness, i.e. , did not result well in agreement with the experimental results, with the theoretical values always exceeding the experimental ones. This could be due to several reasons, for example the introduction of the input values required by the models (elastic constant, coefficient of thermal expansion and fiber strength) [36,37] could not be coherent with the actual values of the composite after sintering. Increase of residual stresses or, most probably, decrease of tensile fiber strength due to chemical alteration upon thermal treatments could occur. This aspect deserves surely further investigation and, in particular, the actual properties of the fibers in the ceramic should be measured. In the case of HfC a modest increase was recorded, however, this was attributed more to the increase in the final density of the matrix, which in turn determined an increase of its toughness, rather than effective toughening mechanisms activated by the fibers. This hypothesis is corroborated by previous results obtained on dense HfC materials, for which the fracture toughness without any reinforcement was 3.6 MPa m1/2 [10]. Flexural strength is another property affected by the fiber addition. It should be noticed that for TaC-based matrix, room temperature strength data are quite scattered. Critical defects can be large TaSi2 pockets, especially when micro-cracked, as observed in Fig. 3c. Moreover, being TaC a very hard and brittle ceramic, surface defects due to machining may also affect the strength. A notable data scattering for TaC was reported also by Zhang et al. [9]. With the addition of fibers, the strength of this hard carbide did not statistically change. A slight decrease of the average values strength was observed, accompanied by a decrease of the standard deviation. Very  \\x0c', '1374  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  Fig. 7. External surface (top) and cross section (bottom) of (a, c) TCTf and (b, d) HCTf after the oxidation test at 1600 1C for 5 min in air. The insets in (a and b)  show the appearance of  the specimens upon oxidation.  likely, the dimension of the critical defects generated by fibers was not significantly different from pristine ones. On the contrary, for HfC, matrix strength values are rather low, due to incomplete sintering. Indeed, typical values measured on similar unreinforced matrices range from 460 to 540 MPa for hot pressed or pressureless sintered matrices with TaSi2 or MoSi2 as sintering agents [14,20,31]. With the addition of fibers, a slight improvement was detected which is due to the increase of final density. In this case, SiC fibers acted more as a sintering aid rather than a strengthening phase. This is clearly  demonstrated by microstructural analyses; indeed large porous regions with hafnia particles observed in the matrix (HCT-2) were replaced by dense mixed SiC-HfO2 regions which were just localized in the proximity of SiC fibers, as shown in Fig. 2b and c. For the same reason, an increase of hardness was observed in HCTf compared to HCT-2 (Table 2). Other properties, such as thermal conductivity, are expected to be affected by the addition of SiC fibers. Although we did not measure the thermal conductivity of the respective matrices, unpublished KTH measured on TaC and HfC matrices with MoSi2  \\x0c', 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1375  as sintering aid showed the same trend as the composites object of this publication, i.e. thermal conductivity increase with the temperature and higher values for TaC-based compounds. For those analogous ceramics with MoSi2 the conductivity was similar to the reinforced composites up to 500 1C, but at 1900 1C the unreinforced composites had a conductivity up to 10 W/m K, higher than the SiC-fiber reinforced ones, since at high temperature, the thermal conductivity of HfC and TaC increases, whilst it decreases for SiC. Discussing the effects of SiC fiber on the overall thermal conductivity is not straightforward because we do not thermal conductivity up to 1900 1C, but we know the fiber can assume that the introduction of SiC fiber in TaC or HfC matrix causes a conductivity decrease at high temperature. Moreover, microstructural changes due to exposition at high temperature, such as SiC active oxidation in the subsurface layer resulting in the creation of voids could also affect this property. HCTf section after thermal conductivity test is shown in Fig. 9a and subsurface voids are pointed out in the inset. Fig. 9b and c show the surface covered by SiC due to the graphite-based painting used for the test. However we can assess that the lower values for HCTf are due both to the higher porosity of its matrix (about 7%) as compared to the dense TCTf and to the intrinsic lower conductivity of HfC.  As for the oxidation behavior, from SEM analyses reported in Fig. 7 it is clear that, both for TaCand HfC-composites, 15 vol% of fiber were not enough to form a continuous silica layer and therefore to form an efficient barrier against oxygen penetration at 1600 1C. There can be several reasons for such behavior: there is not enough silica to fill all the volume expansion associated with the formation of the oxides, or the thermodynamic involved causes volatilization of silica to gaseous SiO on the surface. Though, we can confirm that the introduction of SiC fiber improved the oxidation resistance of TaC, as evidenced in Fig. 5. In summary, the addition of SiC fibers did not produce any measurable benefit in terms of fracture toughness or strength. Compared to available literature data on composites HfC and TaC reinforced with SiC particles, it can be observed that TaC is reinforced more efficiently by SiC particles than SiC fibers. For instance, in the work of Liu et al. [38], addition of 10 vol% SiC particles resulted in a 3-pt bending strength of 660 MPa, and fracture toughness approaching 6 MPa m1/2. On the contrary for HfC, addition of SiC particles resulted in toughness values of 2.2-2.8 MPa m1/2 (i.e. even lower than upon addition of fibers) and similar 4-pt bending strength values (396 MPa) [39]. A significant contribution from SiC fiber mainly invested the densification behavior, especially in the case of HfC. If on one  Fig. 8. Crack path crossing the fibers generated by a 98.1 N indentation onto TaC (a) and HfC (b) with fibers and details of  the advancing crack along the  fiber/matrix boundary.  Fig. 9. HCTf after thermal conductivity test: (a) section and detail of the voids in the subsurface, (b) SiC grains on the surface, and (c) detail of the surface showing  HfC grains among SiC ones.  \\x0c', 'side the fiber released some silica-based phases that were beneficial for obtaining a denser matrix, on the other side a too strong interaction occurred between the fiber and matrix causing notable fibers modification of morphology and properties and not allowing efficient toughening mechanisms.  5. Conclusions  TaCand HfC-based composites containing 15 vol% short SiC Tyranno SA3 fibers were produced by hot pressing at 1700-1750 1C. TaSi2 was selected as sintering additive to lower the maximum sintering temperature and preserve SiC fibers. The addition of SiC fibers enabled densification of carbides at temperatures even lower than the respective matrices, but no benefits in terms of fracture toughness or strength were observed. 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Bubsey, Fracture toughness calcula tion  from maximum load  in  four  point  bend  tests  of  chevron  notch  specimens,  Int. J. Fract. 16 (1980) R137-R141.  [31] D. Sciti, L. Silvestroni, S. Guicciardi, D. Dalle Fabbriche, A. Bellosi,  Processing, mechanical properties  and oxidation behavior of TaC and  HfC composites containing 15 vol% TaSi2 or MoSi2, J. Mater. Res. 24 (6) (2009) 2056-2065.  [32] C.K. Jun, P.T.B. Shaffer, Thermal expansion of niobium carbide, hafnium  carbide and tantalum carbide at high temperatures, J. Less Common Met.,  Elsevier Sequoia S.A., Lausanne - Printed in The Netherlands, 323, Short  Communication, Research and Development Division, The Carborundum  Company, Niagara Falls, N. Y. 14302 (U.S.A.)  (received December 19th,  1970).  [33] R.S. Roth, J.L. Waring, Diagram 4448: Ta2O5-SiO2, for Ceramists, 1970.  in: Phase Diagrams  L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1376  \\x0c', 'L. Pienti et al.  / Ceramics International 41 (2015) 1367-1377  1377  [34] L. Charpentier, M. Balat-Pichelin, D. Sciti, L. Silvestroni, High tempera [37] F.F.  Lange,  The  interaction  of  a  crack  front with  a  second-phase  ture oxidation of Zrand Hf-carbides:  influence of matrix and sintering  dispersion, Phil. Mag. 22 (1970) 983-992.  additive, J. Eur. Ceram. Soc. 33 (2013) 2867-2878.  [38] L. Liu,  F. Ye, Z. Zhang, Y. Zhou, Microstructure  and mechanical  [35] S. Guicciardi, L. Silvestroni, M. Nygren, D. Sciti, Microstructure  and  properties of  the spark plasma sintered TaC/SiC composites, Mater. Sci.  toughening mechanisms  sintered ZrB2 forced by SiC whiskers or SiC chopped fibers, J. Am. Ceram. Soc. 93  in spark plasma  ceramics  rein (2010) 2384-2391.  Eng. A529 (2011) 479-484.  [39]  J.X.  Liu, X. Huang, G.J.  Zhang,  Pressureless  sintering  of  hafnium  carbide-silicon  carbide  ceramics,  J. Am. Ceram. Soc.  96  (6)  (2013)  [36] M.  Taya,  S. Hayashi, A.S. Kobayashi, H.S. Yoon,  Toughening  of  a  1751-1756.  particulate-reinforced ceramic-matrix composite by thermal  residual  stress,  J. Am. Ceram. Soc. 73 (1990) 1382-1391.  \\x0c']"
},{
  "_id": 131,
  "PDF": "Modeling Oxidation Kinetics of SiC-Containing Refractory Diborides.pdf",
  "Text": "['Modeling Oxidation Kinetics of SiC-Containing Refractory Diborides  T. A. Parthasarathy,  ‡,§,†  R. A. Rapp,  ¶  M. Opeka,k and M. K. Cinibulk  ‡  ‡  Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright-Patterson AFB, Ohio 45433-7817  §  UES, Inc., Dayton, Ohio 45432  ¶  The Ohio State University, Columbus, Ohio 43235  kNaval Surface Warfare Center, Carderock, Maryland 20817  Experimental data on the oxidation kinetics of SiC-containing 1473-  diborides  of Zr  and Hf  in  the  temperature  regime  of  2273 K are interpreted using a mechanistic model. The model  encompasses  counter-current gas diffusion in the  internal SiC  depleted  zone,  oxygen  permeation  through  borosilicate  glass  channels in the oxide scale, and boundary layer evaporation at  the surface. The model uses available viscosity,  thermodynamic  and kinetic data for boria,  silica, and borosilicate glasses, and  a  logarithmic mean  approximation  for  compositional  varia tions. The  internal  depletion  region  of SiC is modeled with  CO/CO2 counter diffusion as the oxygen transport mechanism. Data reported for pure SiC in air/oxygen, for ZrB2 containing varying volume fractions of SiC, and for SiC-HfB2 ultra-high temperature ceramics (UHTCs) by different investigations were  compared  with  quantitative  predictions  of  the model.  The  model  is  found  to  provide  good  correspondence with  labora tory-furnace-based  experimental  data  for weight  gain,  scale  thicknesses, and depletion layer thicknesses. Experimental data  obtained from arc-jet  tests at high enthalpies are found to fall  well outside the model predictions, whereas lower enthalpy data  were  closer  to model  predictions,  suggesting  a  transition  in  mechanism in the arc-jet environment.  I.  Introduction  R EFRACTORY  diborides,  especially ZrB2 and HfB2, with being studied with great interest  SiC additions  are  because  of  their  high  thermal  conductivity,  high melting  point, and moderate resistance to environmental degradation temperatures.1-3 This these two-phase composites makes  in  air  at  very  high  combination  of  properties of  them prom ising for use  in hypersonic vehicles as  leading edge fluxes.1  compo nents,  which  are  subject  to  high  heat  Several  additions  and compositional  refinements  are being  tried to  enhance  their oxidation resistance, but SiC-containing ZrB2 or HfB2 are the most studied, with sufficient data to enable a modeling-based analysis and understanding of their oxidation behavior.4-19 A brief summary of experimental observations lows.4-19 The oxidation behavior kinetics despite formation of a porous oxide of  is  as  fol is dominated by parabolic  the refractory  metal. This protective behavior  is attributed to the borosili cate glass  that fills  the pores of  the porous oxide  scale and  often  also  results  in  a  glassy  external  layer,  as  seen  from  Fig 1(a)  (courtesy Prof. E Opila, University of Virgina). In a  furnace under static air, a continuous external glassy layer is  formed, but at high temperatures,  it begins  to flow due to a  decrease in viscosity giving rise to some scatter in experimen tal data. The underlying porous oxide is almost always filled  with  the  borosilicate  glass. Underneath  this  oxide  layer,  some  investigations have  reported a depleted region where  SiC is absent but  the diboride is  still  intact as,  for example,  seen in Fig. 1(a). The depleted region is absent at  low tem peratures, but has been reported in samples exposed to high temperature  air. However,  there  is  an inconsistency  in the  reported  data with  some  studies  not  finding  any  depleted  layers under the same conditions where others have observed  them.7,10,14,20,21  In some work, a partially consisting of SiC and a Si-O-C phase,  depleted  region  is reported. The high est  temperatures are achieved under arc-jet conditions, where  typically very thick oxide layers with a very thin or no exter nal glassy layer are reported for very short times. A depletion  zone has  also been observed under  these  conditions. Some  works have suggested that convection currents can be significant during oxidation.17-19 To obtain a quantitative model that interprets these data was the objective of this work.  In an initial  study, we reported on a model  for  the oxida tion of monolithic diborides of Zr and Hf in the temperature region of 1073-2673 K.22-24 In the present work, we add the complexities arising from the presence of SiC and its oxida tion products,  silica and CO. We retain all  the essential ele ments of  the prior model,  including volumetric  effects  from  phase  change of the oxide, and transitions in mechanistic temperature is increased.24 As in the prior model,  regimes as  the  refractory oxide  in the  scale  is assumed to be  imperme able to oxygen due to tivity.22 The effects of  the  known  low electronic  conduc the presence of SiC are the formation  of  a borosilicate  glass  instead of  a boria  glass with all  its  consequences,  the  formation and  transport of  gaseous CO  and SiO, and the evaporation of SiO and SiO2 at The properties of silica and boria are known but modeling  the surface.  their  variation with  composition  needs  interpolation. The  vapor pressures of B2O3 and SiO2 are known as a function of temperature, and that of SiO is known as a function of  temperature and oxygen partial pressure.  The model  is  able  to interpret most of  the  experimental  data reported on ultra-high temperature  ceramics  (UHTCs).  Weight gain, oxide thickness, external glass thickness, and internal depletion layer thickness for both SiC-ZrB2 and SiC -HfB2 are found to be in reasonable agreement with experimental data, but with some exceptions mostly related to  arc-jet  test data. The equations that constitute the model are  presented first;  this is followed by a section where the predic tions  are  compared with  the  available  experimental  data.  The final  section discusses  the merits and limitations of  the  N. Jacobson—contributing editor  Manuscript No. 29879. Received June 15, 2011; approved October 03, 2011.  This work was  supported by USAF Contract # FA8650-10-D-5226 which included  funding  from US Air Force Office  of  Scientific Research  (AFOSR), monitored  by  Dr. Ali Sayir.  †  Author to whom correspondence should be addressed. e-mail:  triplicane.parthasar athy@wpafb.af.mil  338  J. Am. Ceram. Soc., 95 [1] 338-349 (2012)  DOI: 10.1111/j.1551-2916.2011.04927.x  © 2011 The American Ceramic Society  Journal  \\x0c', 'model,  including  possible  reasons  for  discrepancies  and  suggestions for future work.  II.  The Model  (1)  Model Framework and Key Assumptions  The morphology of  the oxidation scale assumed in the model  was derived from published cross-sectional microstructures of oxidized UHTCs.4,5,8,10,13,21,25 The key elements of  the model  are  shown  in Fig. 1. Figure 1(a) shows a sample microformed on a ZrB2-SiC material. The in Figs. 1(b) and (c), details the assumed mor structure of  the  scale  illustration,  phology of  the oxidation products,  representing steady-state  conditions. The substrate is a composite of SiC and MeB2 (Me = Zr, Hf) with fs being the volume fraction of SiC. The oxidation product, viz. the scale, consists of disconnected (in  cross-section) MeO2 grains with a continuous porous of solid) region that is filled with liquid borosilicate.  (void  Throughout  the manuscript, pore  refers  to a region void of  MeO2 but filled with glass, except at very high temperatures when glass evaporates leaving pores behind [Fig. 1(c)]. The  MeO2 grains are taken to be impermeable to oxygen, as explained in our prior work.22 Molecular oxygen from the  ambient  dissolves  in  borosilicate  as molecular O2, the external glassy phase, and permeates through the  diffuses  across  porous channels of glass, eventually reaching the interface i2,  where  it  reacts with the diboride. A portion of the oxygen (= volume)  flux, assumed to be proportional  to areal  fraction  of  SiC,  is  transported within  the  depleted  region  (void  of  solid but filled with gases)  to oxidize SiC through a medium  of a gas mixture of CO and CO2, similar to the model by Holcomb and St. Pierre for HfC oxidation.26 The  gaseous  reaction product, CO, of  the SiC oxidation is  assumed to  either diffuse or bubble through the glassy scale (as observed in ref. 17-19) and is assumed not  to be rate limiting, consisSiC.27 Fig tent with  prior  assumptions  for  oxidation  of  ure 1(b) depicts  the expected scenario at higher temperatures  (likely above 2000 K) where  evaporation of B2O3 and SiO2 sufficiently high that an external  from the external  surface is  glassy layer cannot be supported and the glassy layer recedes  inwards.  In all  regimes,  the diffusivity of gases  in a multi component gas mixture and the  effect of Knudsen diffusion  [when an unfilled pore was present as in Fig. 1(b)] were modeled as detailed in prior work.22  The volume fraction of SiC in the substrate is  fs, and the volume fraction of porosity (filled with glass) within the MeO2 region is fp. As in prior work, the pore fraction, fp, (filled with glass) is taken to change as the temperature crosses the monoclinic-tetragonal  phase  transformation  temperature,  Ttrans.  The parameters  fs and fp are related to the volume fraction,  (a)  (b)  (c)  1  2  3  a  )  (  )  (  )  (  )  (  )  (  )  (  2  2  2  3  2  3  2  g  SiO  l  SiO  g  SiO  l  SiO  g  OB  lOB  fs  CO (assumed Not rate limi(cid:415)ng)  l12  l23  l3a  1 fs  SiO2(g) SiO(g) B2O3(g)  Boundary Layer diffusion  CO  SiO  CO  SiC  3  2  2  2  2  2  2  2  2  1 2  CO  O  CO  SiO  O  SiO  i1  i2  i3  a  MeB2  SiO/CO,CO2  B2O3 SiO2  B  S  i  O  2  MeO2  SiC  3  2  2  2  2  5 2  OB  MeO  O  MeB  O2 (dissolved)  Evapora(cid:415)on limited by Diffusion through Porous channe ls SiO2(g) SiO(g) B2O3(g)  MeB2  SiO/CO,CO2 B2O3 SiO2  MeO2  SiC  1  2  3i  a  i1  i2  i3i  a  O2 (diss.)  Fig. 1.  (a) SEM image of  the microstructure of oxidation scale  formed on a ZrB2-SiC sample the oxidation products and morphology assumed in the model. At lower temperatures  (Courtesy: E Opila, Univ. Virginia)  (b,  c)  Schematic  sketches of  (b)  external glassy scale  forms,  whereas at higher temperatures or high ambient flow (c), the glassy scale recedes inward due to evaporative loss of SiO2 and B2O3.  January 2012  Oxidation Modeling of SiC-Containing Refractory Diborides  339  \\x0c', 'fMeO2 , of MeO2 in the glass + MeO2 region (2-3), and fraction of glass, fg, (borosilicate) in the scale as follows.  fp ¼ f1 for T  \\x0c', 'One of  them, PCO-i1, implicit equation. The only unknowns in the above equation region 2-3 (Fig. 1), l12. The flux equations that apply  requires numerical  solution of  the  set are  the oxygen flux,  JO2\\x0032  j  j,  in the  and the length of  the depleted zone,  JO2 \\x0032  j  j  is  obtained from solving 2-3, discussed below.  the  to  region  (3)  MeO2 + B2O3-SiO2 Glass Region (2-3)  The  equations  that  govern  the  oxygen  flux  in  region  2-3,  between interfaces i2 and i3, are given below.  JO2\\x0032  j  j¼PO2\\x00B2 O3\\x00SiO2  PO2\\x00i3\\x00PO2\\x00i2  l23 PO2\\x00a\\x00PO2\\x00i3  fg  JO2\\x00a3  j  j¼PO2\\x00B2 O3\\x00SiO2 ð3aÞ  l3a  Since; JO2\\x00a3  j  j¼ JO2 \\x0032  j  j  PO2 \\x00i3¼l23PO2 \\x00B2O3 \\x00SiO2 ð3aÞPO2 \\x00aþfg l3aPO2\\x00B2O3\\x00SiO2 PO2 \\x00i2 l23PO2\\x00B2 O3\\x00SiO2 ð3aÞþfg l3aPO2\\x00B2O3\\x00SiO2  (7)  Here, Π refers  to the permeability  (rate  of  transport  in  moles per unit area per unit  time per unit partial pressure  gradient). The permeability in region 2-3, as  in region  3-a  is  different  from  that  region 3-a is  likely to be a boria-defi cient  region resulting from faster  evaporation of boria than  silica.  The unknowns  that  remain are  the  three  lengths,  l12, l23, the three regions. These are allowed to evolve with  and l3a of time from an initially small value within a numerical simula tion model. The evolution equations for l12 and derived from flux and molar volumes are given below. The rate of change  of  l23  is given by the number of moles of MeO2 unit area per unit time multiplied by the molar volume of  formed per  MeO2. The rate of formation of MeO2 flux of oxygen in through region 3-2, less than that used for  is given by the total  oxidizing SiO(g)  to SiO2 and CO(g) to CO2, which is the SiO(g) flux in region 1-2 as given in Eq. 4. The  3/2  times  recession rate of MeB2 and SiC are simply related to the rate of change of through molar volumes and volume frac l23  tions. Finally,  the depleted zone  length change  is  given by  the difference in recession between MeB2 and SiC. Thus, we obtain the following equations.  dl23  dt  ¼ VMeO2 ð JO2\\x0032  j  j \\x00 3  2  JSiO\\x0012  j  jÞ 2  5  1  fMeO2  dRMeB2  dt  ¼ dl23 dt  VMeB2 fMeO2  VMeO2 ð1 \\x00 fs Þ  dRSiC  dt  ¼ JSiO\\x0012  j  j VSiC  fs  l12 ¼ RSiC \\x00 RMeB2  (8)  In the above set of equations, t  is time, R is recession, V is  molar  volume,  and J  is flux  (moles per unit  area per unit  time).  (4)  External B2O3-SiO2 Glass Region (3-a)  The evolution with time for the thickness,  l3a, of is given by the rate of production of boria  the external  glassy layer 3-a,  and silica less  the rate of  loss by evaporation at the surface in region 2-3. From the  and the amount occupied by glass  literature, B2O3(g), SiO(g), and SiO2(g) are known to be the dominant gaseous species in the B2O3-SiO2 system in the  temperature and oxygen partial pressures of  interest  (see for  example, ref. 7).  dl3a  dt  ¼  dl23  dt  \\x0c\\x0c \\x00 JSiO\\x00vap fMeO2 \\x00 JB2O3 \\x00vap VMeO2 \\x00 JSiO2 \\x00vap \\x0c\\x0c ¼ Dspecies  \\x12  \\x13  VB2 O3 þ ð JSiO\\x0012  j  j  \\x0c\\x0cÞVSiO2 \\x00 dl23 dt  fg  Jspecies\\x00vap  RT  105 Pspecies\\x00vap dbdry  ; dbdry  ¼ 3  2  s ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi  lspecimen  vfluid  gflluid qfluid  \\x14  \\x15  1=6 ðDspecies Þ1=3  (9)  In  the  above  equation  set,  the  subscript ‘vap’ refers to the specimen, g, ρ, and v refer to viscosity, density, and velocity of  evaporation,  lspecimen is  the length of  the ambient fluid. d) The last equation gives the boundary dependent evaporation of species.30  layer  (thickness  The viscosity of  the glassy layer could limit  the maximum  thickness,  l3a-max, of  the external layer that can be supported under gravitational forces. As in the previous work,24,31 from falling liquid films,31 under  the  theory of  laminar flow,  this  limiting thickness is given by:  l3a\\x00max ¼  3ðMB2O3\\x00SiO2  CB2O3\\x00SiO2 lspec ÞgB2O3\\x00SiO2  gq2  B2O3\\x00SiO2  sinð/Þ  \"  #  1=3  CB2 O3\\x00SiO2 ¼ dl3a  dt  1  VB2 O3\\x00SiO2  (10)  where  subscript B2O3-SiO2 at the surface region (3-a), M refers to molecular weight, g the acceleration due to gravity, u the orientation of to gravity, and CB2O3\\x00SiO2 at which the borosilicate is added to the surface in moles per  refers  to the borosilicate present  the sam ple with respect  refers  to the rate  unit area per unit  time.  Finally,  the  net weight  gained, weight  of  oxygen  con sumed, and weight evaporated can be given as:  Wg ¼ l23 ðfMeO2 qMeO2 þ fgqg Þ þ l3aqg \\x00 RSiC fs qSiC \\x00 RMeB2 ð1 \\x00 fs ÞqMeB2  WO2 ¼ 5 2  RMeB2  ð1 \\x00 fs Þ  VMeB2  MO2 þ 3 2  RSiC  fs  VSiC  MO2  Wevap ¼  X  ðSiO2 ;SiO;B2 O3 Þ  Mspecies  Z  t  0  Jspeciesdt þ RSiC fs VSiC  MCO  (11)  The  variation in time of  the oxide  scale  thicknesses  and  the  various weight gain/loss can be computed numerically (1)-(11) using an evolutionary algorithm which is  using Eqs.  started with a choice of,  initially, arbitrarily small value  for  l12,  l23, and l3a.  (5)  High-Temperature Regime  At high temperatures (likely above 2000 K) and/or at high flow rates of ambient fluid (>1 m/s), where rates of SiO2(g) and/or SiO(g) are sufficiently high, nal glassy layer will be lost and the glassy liquid region will recede into the region 2-3, as shown in Fig. 1(b). Using 3i  the  evaporation  the exter to  identify the location of  the receding glassy layer,  the follow ing equations describe the oxygen flux balance which fixes the partial pressure of oxygen, PO2 \\x00i3i , at the location 3i,  the interface i3i at  January 2012  Oxidation Modeling of SiC-Containing Refractory Diborides  341  \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c\\x0c \\x0c', 'The diffusive fluxes of  gaseous  species  (SiO2, B2O3, the interface i3i. Tak and  SiO) determine the rate of recession of  ing the ratio of the depth of region 2-3 as q,  the glassy region to the length  l23 of as:  the evolution equation for q is given  Jspecies\\x00vap ¼ Dspecies\\x003ia  RT  105 Pspecies\\x00vap  ð1 \\x00 qÞl23  dq  dt  ¼ 1  l23 \\x00 JSiO2\\x00vapVSiO2 \\x00 JB2 O3\\x00vapVB2O3 g \\x00 q  \\x14  1  fg  fJSiO\\x0012VSiO2 þ ðJO2\\x003i2 \\x00 3  2  JSiO\\x0012 Þ 2  5  VB2O3  dl23  dt  \\x15  (13)  The  right-hand side of  the  evolution equation is  the  rate  of production of B2O3 and SiO2 less the evaporative flux; last term accounts for the increase in during the time  the  l23  increment. The weight gain of  the sample is given as:  Wg ¼ ql23 ðfMeO2 qMeO2 þ fgqg Þ \\x00 Rs fs qSiC \\x00 RMeB2 ð1 \\x00 fs ÞqMeB2  (14)  Again,  the variation in time of  the oxide  scale  thickness  and the weight  gain/loss  can be  computed using  an evolu tionary algorithm.  III.  Model Predictions and Validation  A numerical  code written in Fortran was used for  all  the  model predictions shown in this work. The temperature-time  input  variables  are  the  history,  environment  parameters  (total  pressure,  fraction  of  oxygen),  specimen  orientation,  length, fluid velocity, volume fraction of SiC, size of SiC par ticles. All  the  thermodynamic and kinetic parameters  taken  from the literature reside within the code;  these include oxy gen permeation and viscosity of  silica and boria, equilibrium  constants  for  all  the  reactions,  and vapor pressures of  the  species. An infinitesimal  layer of oxide  and  a  glassy  layer  were assumed present at an infinitesimal  start  time and the  time evolution of all the parameters were calculated as dictated by the equations. An arbitrary time-temperature profile  could  be  thus  simulated  and  the  resulting  thicknesses  of  oxide,  glassy  layer,  and depleted layer  could be  computed.  From these,  the recession,  the weight gain, weight of oxygen  consumed, and weight of  evaporated species  could be  com puted.  (1)  Parameters  A list of variables used in the model  is shown in Table I with  a brief description. The model uses data from the  literature  for all  thermodynamic quantities, viz.  equilibrium constants  and vapor pressures of compendium by Barin.32  species, which are  available  in the  (A)  Gas Diffusivity:  The diffusivities of gases in a mul tigas solution were calculated using parameters given by Svehla,33 and the methods outlined in Ref. 22 The gas diffu sivity applied to the depleted region as well as  the boundary  layer  diffusion  region  at  the  surface  for  evaporation. The  Knudsen effect on diffusivity was  included as determined by 0.5 lm in  the  pore  radius,  assumed  to  be  all  predictions  shown herein;  this choice was based on reported microstruc tures  in  the  literature. The Knudsen  effect  only  comes  in  when the pores dry out under conditions where the evapora tion is  sufficiently fast  [Fig. 1(c)]. The Knudsen effect  is also  calculated for the depleted zone as determined by SiC size.  (B)  Boria Activity:  The calculation of equilibrium parthe ZrB2-ZrO2 the partial pressures of volatile species at  tial pressure of oxygen at  interface  (i2) and  the outer  surface  require the activity of boria in the borosilicate glass. The activity isotherm in the boria-silica system was measured at  1475 K and is found to be nearly ideal, although there is a slight deviation from ideality.28 The B2O3-SiO2 system is also known for the lack of any phase separation.34 Based on  these, a unit activity coefficient was assumed in this work.  (C)  Glass Viscosity:  The  temperature-dependent  vis cosity of boria was obtained from the works of Eppler and Li et al.35,36 The data on viscosity of silica were reviewed by Doremus,37 and as per his conclusion, the data of Urbain et al.38 were used for the temperature range 1400°C-2500°C and that of Hetherington et al.39 for the temperature range 1400°C-1000°C. For  compositional dependence of  viscosity  on boria content, a log-mean interpolation scheme was used consistent with semiempirical models of glass viscosities.40,41  (D)  Oxygen Permeability:  The permeabilities of oxy gen in liquid boria and silica were obtained by fitting to data  from several sources. The boria data were obtained from Tokuda et al.,42 Luthra,43 and Schlichting.44 The oxygen per meability in silica was obtained from the works of Lamkin et al.45 and Courtright.46 For  the compositional dependence,  a log-mean approximation was used for interpolation consistent with what might be expected from Stokes-Einstein relation as suggested by Karlsdottir and Halloran.17  The boria concentration at the surface layer must be lower in the interior oxide + glass region due to evaporative losses of boria being higher than silica. The surface layer  than that  concentration depends on the evaporative rates, diffusion of  B in B-SiO2, and the degree of convective mixing. This calculation is beyond the scope of this work. Further, the authors  have found no experimental data on this. Hence,  the surface  boria concentration was assumed to be a constant fraction of ~0.72 (interior B2O3 this was the only loose parameter in the model and valconcentration of ~0.58 and ~0.65)  the  interior  concentration is  for  20%  SiC);  ues of 0.8 and 0.9 (B2O3 for HfB2 and ZrB2, respectively, dence. The effective pore fractions, fp, for zirconia and hafnia were taken to be the same as was in ref. 24 for monolithic  gave  the  best  correspon diboride oxidation (0.03 and 0.04 for HfO2 and ZrO2, respectively). The aliovalent dopant concentration, in  Cdopant,  MeO2 was taken to be less than 100 ppm, which permits the neglect of oxygen permeation through the MeO2 phase (an excellent assumption) (see ref. 22-24). The ambient fluid  velocity was obtained from the  literature when reported;  it  varied from 0.0001 to 150 m/s, with the smaller values corre sponding to static air, and the higher values close to arc-jet  conditions  (behind the shock wave). The effect of fluid flow  on  evaporation  rates was  included  in  the model  using  a  JO2\\x003i2 ¼ PO2 \\x00B2O3 \\x00SiO2 105 PO2\\x00a \\x00 PO2\\x00i3i  PO2\\x00i3i \\x00 PO2 \\x00i2  ql23  fg  JO2\\x00a3i ¼ DO2\\x003ia  RT  ð1 \\x00 qÞl23  fg  Since JO2\\x00a3i ¼ JO2\\x003i2  PO2 \\x00i3i ¼ qRTPO2 \\x00B2O3\\x00SiO2 PO2 \\x00i2 \\x00 105 qDO2\\x003iaPO2\\x00a \\x00 RTPO2\\x00B2O3\\x00SiO2 PO2\\x00i2 qRTPO2\\x00B2O3 \\x00SiO2 \\x00 105 qDO2\\x003ia \\x00 RTPO2\\x00B2 O3\\x00SiO2  (12)  342  Journal of  the American Ceramic Society—Parthasarathy et al.  Vol. 95, No. 1  \\x0c', 'boundary layer calculation (as detailed in ref. 22-24), but  the  possible  thinning  of  the  external  glassy  layer  from shear  forces was not  included in the model.  (2)  Comparison of Model Predictions with Experimental  Data  (A)  SiC:  There are a lot of oxidation data for SiC in  the  literature,  and  comparing  the model  to  these  data  is  important as a first check. As the oxygen permeability in sil ica was used in the model,  the comparison of model  to data  must be taken as validation of  the assumption that  the kinet ics  of  oxidation  is  limited  by  oxygen  permeability  in  the  glassy phase,  at  least  for SiC. Figure 2  shows  a plot  that  compares  the  available  data  for  single  and polycrystalline investigations47-53 in the data. Figure 2  SiC in various  forms  reported by different  in pure oxygen;  there is a large scatter  includes data collected on high purity SiC by Ramberg et al.,47 and Ogbuji and Opila53 as well as data obtained by a Costello and Tressler49 on SiC made by different processing  methods. The model  is  seen to fit  fairly well with the poly crystal data, and the fast oxidizing (Si  face) data for the sin gle crystal.  (B)  SiC-ZrB2:  A comprehensive work on the effect of  SiC content on the oxidation kinetics of UHTCs was performed by Talmy.54 The  study was  conducted using a ther mogravimetric apparatus  in flowing air. The weight gain of  samples exposed for 2 h in air et al.15 have  is  reported.  In a more recent  work, Wang  conducted  a  similar  study using  SiC volume fractions of 0, 5, 10, 15, and 20 vol% at atures up to1600°C for 1h in an air  temper furnace. Figure 3 com pares the data with the model prediction for weight gain as a  function of SiC volume percent. Figure 3(a) shows data from Talmy54 et al.15 The  and Fig. 3(b)  shows data  from Wang  model predictions  for  the  sample weight gain and the  total  weight of oxygen consumed are  shown as  solid and dashed  lines. The sample weight gain is the weight of  the sample and  oxidation products  less  the weight of  external glassy phase  predicted  to  be  lost  by  evaporation  or  viscous  flow. The  weight of oxygen consumed will be the total weight change if  all of  the oxidation products were  retained on the sample. 1200°C from Wang et al. The model agrees with the trends very well.  Figure 3(c)  includes  time dependence data  at  The temperature dependence of weight gain and scale thickness data on a fixed composition of 20 vol%SiC-ZrB2  1.E+01 5.50E-04  1.E+02  1.E+03  1.E+04  6.00E-04  6.50E-04  7.00E-04  P  a  r  a  o b  i l  c  r  a  t  e  c  n o  s  t  a  n  t  ,  K  p  a  r  ,  n  m  2  /  m  i  n  1 / Temperature, K  S iC  s intered , o xygen : Co stello, Treess ler [50]  S iC HP, o xyg en Co s tello, Treess ler [50]  SC S iC-f as t -o xygenCo stel lo, Treess ler [50]  SC S iC s low -o xyg enCo stello, Treess ler [50]  Ramb erg , Wo rrel [48]  CVD S iC Og b uji, Op ila [54]  CVD S iC : Harris [53]  S i f ace Zheng et al. [52]  C f ace Zheng et al. [52]  SiC  Fig. 2.  The oxidation kinetics, expressed as parabolic rate constant,  in  oxygen  of  single  crystal  and  polycrystal  SiC in  various  forms  reported  by  various  investigators  are  shown  compared with  the  model. The model prediction is  shown as a solid line, whereas  the  experimental data are shown as dotted lines.  Table I.  A List of Symbols Used in This Work, with a Brief Description and Units  Symbol  Units  Description  fp fs  Effective volume fraction of pores in MeO2 that Volume fraction of SiC Volume fraction of MeO2 in the MeO2-glass region (2-3) Volume fraction of glass region in the MeO2-glass region (2-3) Activity of species at interface i2  is permeable to gas  fMeO2  fg  aspecies-i2 RMeB2 RSiC Vspecies Pspecies Jspecies-12 Dspecies-12  m  Recession of MeB2 Recession of SiC  m m3  Molar volumes of species  atm mol/m2-s m2/s  Partial pressure of species Flux of species from interface 1-2 Diffusivity of species in region 1-2  R  J/mol-K  Universal gas constant  PO2 \\x00B2O3\\x00SiO2  mol/m-s-atm  Permeability coefficient of oxygen in liquid borosilicate internal depletion (=Rs \\x00 Rme) Thickness of zirconia region in the scale Thickness of external glassy layer of B2O3-SiO2(l) Thickness of zirconia scale over which B2O3-SiO2(l) Temperature  I12 l23 I3a  m  Depth of  m  m  q  m  is present  T  K  t  s  Time  Mi  kg/mol kg/m3  Molecular weight of species i  ρi gi  Density of species i  Pa-s mol/m2-s mol/m2-s  Viscosity of species i  CB2 O3\\x00SiO2  Rate of addition of boria to the external boria scale  Jspecies-vap dbdry Ispecimen Vfluid  Rate of evaporation of species i at  the external surface  m  Boundary layer thickness for surface evaporation  m  Length of specimen  m/s kg/m2 kg/m2 kg/m2  Velocity of ambient fluid  Wg  Net change in weight per unit area  WO2  Weight of O2 consumed per unit area Weight of evaporated species per unit area  Wevap  January 2012  Oxidation Modeling of SiC-Containing Refractory Diborides  343          \\x0c', 'in static laboratory air have been reported by Carney et al.10  This  study took special care to document  the effect of glass  flow in two ways. They measured the weight changes of both  the  sample  and the  crucible onto which some of  the  glass  had spread. They also measured the variation in scale thick ness with sample  surface orientation,  and found significant  scatter. Further,  they reported on samples prepared by SPS  and hot press. For  comparison with the model,  the average  values of  all  the  samples  for  a  given condition were used.  The experimental data points attributed to Opila and Halbig  in Fig. 4(a) were calculated from the parabolic constants weight gain reported in their work.13 (sample + crucible)  for  In Fig. 4(a),  the  total  weight  gain  is  plotted  along with  the  model predictions,  showing  good correspondence. The plot  also  shows  that  evaporation  is  significant  above  1873 K.  Figure 4(b) + glass)  shows  the average values  for  total  scale  (oxide  thicknesses  and  oxide  thicknesses.  Figure 4(b)  includes scale thicknesses data measured from images reported by Zhang et al.14 The data of Carney were obtained  in an alumina furnace whereas the data of Zhang et al. were  obtained using a zirconia furnace. A clear  internal depletion  layer  was  not  observed/reported  in  the  experiments.  In  contrast, Fahrenholtz  7  reported a depletion layer of 10 lm 1773 K in a ZrB2-30% SiC sample; 0.6 lm for these conditions. For  after  30 min  at  the  model predicts 20% SiC-ZrB2 case, the model predicts a depletion zone ranging from 1.8 lm at 1673 K to 3 lm at 1873 K which is  only  the  closer  to  observations  by Carney  et al. Figure 5  shows  a  comparison of  the model with data for  the time dependence  of weight gain data.5,10 Data  and  scale  thickness  from two  sources  of  for  hot-pressed  samples  and  SPS-processed  samples and the  sources of  the data are distinguished with  different  symbol  shapes. Once  again,  there  is  a  reasonable  correspondence between the data and the model.  (C)  SiC-HfB2:  Oxidation high-density 20 vol%SiC-HfB2 samples have been collected in static air from 1673 to 2273 K by Carney.11  data  on  SPS-processed  In addition,  the statistical variations  from batch to batch and processing  route  (hot-pressed  versus  SPS)  have  been  studied for the Sevener.20  same  composition  at  1773  and  1873 K  by  Figure 6(a)  compares  the  data  for weight  gain with  the  model predictions. The data for weight gain include separate  measurements + sample  for  the  sample  (open symbols)  and crucible  (filled  symbols).  The model  predictions  include  (a)  (b)  (c)  0  0.1  0.2  0.3  0  20  40  60  80  100  W  e  i  h g  t  g  a  i  n  ,  k  g  /  m  2  Volume % SiC  Data from Talmy [55]  0  0.1  0.2  0.3  0  20  40  60  80  100  W  e  i  h g  t  g  a  i  n  ,  k  g  /  m  2  Volume % SiC  1573 K, 2h  Data from Talmy [55]  0  0.1  0.2  0.3  0  20  40  60  80  100  W  e  i  h g  t  g  a  i  n  ,  k  g  /  m  2  Volume % SiC  1473 K, 2h  1773 K, 2h  Data from Talmy [55]  0  0.05  0.1  0.15  0.2  0.25  0.3  0  5  10  15  20  25  W  e  i  h g  t  G  a  i  n  ,  k  g  /  m  2  Volume % SiC  1873 K, 1h  Data from Wang et al. [15]  0  0.05  0.1  0.15  0.2  0.25  0.3  0  5  10  15  20  25  W  e  i  h g  t  G  a  i  n  ,  k  g  /  m  2  Volume % SiC  Data from Wang et al. [15]  0  0.05  0.1  0.15  0.2  0.25  0.3  0  5  10  15  20  W  e  i  h g  t  G  a  i  n  ,  k  g  /  m  2  Volume % SiC  Data from Wang et al. [15]  1473 K, 1h  1673 K, 1h  0  0.01  0.02  0.03  0.04  0.05  0.06  0.07  0  100  200  300  W  e  i  g  h  t  G  a  i  n  ,  g k  /  m  2  Time, min  Data from Wang et al., [15]  1200C, ZrB2-15SiC  Fig. 3.  (a) The effect of volume percent SiC on the oxidation weight gain in 2 h of UHTC (SiC-ZrB2) samples, as measured by Talmy54 shown compared with the prediction of the model for three different temperatures. The solid lines are the predictions that include the flow of external  glassy layer under gravity and evaporation of boria or  silica. The dotted lines  show the predicted weight of oxygen consumed, which is  the  maximum weight gain that  the samples could have suffered from oxidation in the absence of flow/evaporation of model is shown compared with data from a different investigation, by Wang et al.,15 which used a 1-h hold at temperature. In (c), gain as a function of time for a ZrB2-15 vol%SiC sample at 1200°C reported by Wang et al.15 is shown compared with the model.  the glassy layer.  In (b),  the  the weight  344  Journal of  the American Ceramic Society—Parthasarathy et al.  Vol. 95, No. 1                          \\x0c', 'total weight of oxygen consumed,  sample weight gain, and  weight  evaporated. Evaporation is predicted to dominate at  around 2000 K, and the model  is consistent with the drop in  sample weight gain at  this temperature. In general,  the model  captures the trends well except at  the highest  temperature. In  Fig. 6(b),  the scale thicknesses measured are compared with  the model. The microstructures  of  the  oxidation  product  were  complex  at  the  highest  temperatures,  and  there was  ambiguity in the definition of  the depleted zone. Thus,  the  sum of oxide  scale and depletion zone  (open symbols) was  used for  this  comparison. The  total  scale  thicknesses  (filled  symbols)  are also plotted. The  correspondence between the  model and data is seen to be reasonable, except  for  the data  at  the highest  temperature reported, viz. 2173 K, which devi ates  significantly from the model. However,  these data also  deviate  significantly  from the  extrapolation  of  the  data  at  lower  temperatures. On further  examination of  the  sample,  significant contamination of alkali elements (mainly Ca) pos sibly  from the  furnace  or  crucible  (Ca-stabilized  zirconia)  was detected. Enhanced oxidation of SiC due to the presence of impurities is well known (for example, see Ramberg et al.47)  Figure 7 compares the model predictions time for 20 vol%SiC-HfB2 as a function of in furnace air at 1773 and 1873 K, with data reported by Sevener.20 The data  for weight gain  include SPS-processed material and hot-pressed material with  different  initial SiC particle sizes obtained from different pro cessing  routes. The data  reported for  the  combined weight  gain of  sample and crucible were used. The  correspondence  is once again found to be good, although there  is  consider able  scatter  in  the  data. The model  prediction  plotted  in  Fig. 7 includes the total weight of oxygen consumed showing  that evaporative loss becomes significant at 1873 K.  (D)  Arc-Jet Tests on SiC-MeB2:  The expense of con ducting arc-jet  tests has  limited the number of  investigations  and the extent of data available on oxidation kinetics during these tests. However, data are available for a few SiC-MeB2 compositions. Montverde and Savino have tested a hemiof ZrB2-15 vol%SiC under surface temperature reached  spherical  sample  arc-jet condi~2193 K for  tions where  the  (a)   (b)  1.E 02  1.E 01  1.E+00  4.5E 04  5.5E 04  6.5E 04  W  e  i  g  h  t  g  a  i  n  ,  g k  /  m  2  1/ Temperature, K  Carney e t al. [10]  SPS  Carney e t al. [10]  HP  Opila, Halbig [13]  Mode l Wt evap  Mode l Wt Gain  Mode l Wt O2  2.5 h, S(cid:415)ll Air  Wt O2  Wt gain  Wt evap  1.E 06  1.E 05  1.E 04  1.E 03  4.5E 04  5.5E 04  6.5E 04  S  c  a  l  e  T  h  i  c  k  n  e  s s  ,  m  1/Temperature, K  Ox ide : Carney [10]  (ave HP,SPS)  Total : Carney [10] (Ave HP, SPS)  Ox ide : Zhang et al. [14]  Total : Zhang e t al. [14]  2000  1818  1666  1538 K  Total Scale (oxide+glass)  Oxide Scale  deple(cid:415)on  2.5 h, S(cid:415)ll Air  Fig. 4.  (a) Oxidation weight  gain  in ZrB2-20 vol%SiC samples, et al.10 of exposure, by Carney  measured  in  static  air  after  2.5 h  shown compared to the model predictions. The data from Opila and Halbig13 are also shown. The  solid lines are model predictions  for  weight gain, weight of oxygen consumed, and weight evaporated.  (b)  Average  of  reported values for (oxide + glass)  the  oxide  scale  thickness  (open  symbols) and total  scale thickness  (filled symbol)  for  the  same  exposures;  values measured  on  different  sides  of  the  samples were averaged for this plot. Values from the work of Zhang et al.14 which used a zirconia furnace, are also included. The  solid  lines  are  predictions  for  total  scale,  oxide  scale,  and  depletion  thicknesses. No depletion thickness was reported, depletion was noticed under FIB/TEM investigation.10  but  partial  (a)   (b)   0  0.05  0.1  0.15  0.2  0  50  100  150  200  250  300  W  e  i  h g  t  g  a  i  n  ,  k  g  /  m  2  Time (min)  1873K-Carney et al.[10] SPS  1873K-Carney et al.[10] HP  1873K Model  1600K Model  1900K Lev ine et al.[5]  1600K Lev ine et al.[5]  0.0E+00  2.0E-05  4.0E-05  6.0E-05  8.0E-05  1.0E-04  1.2E-04  1.4E-04  0  100  200  300  S  a c  l  e  t  h  i  k c  n  s s e  ,  m  Time, min  1873K Ox ide-Carney et al.[10] HP 1900K ox ide Lev ine et al.[5] 1873K Ox ide+glass -Carney et al.[10]HP 1873K ox ide+glassCarney et al.[10]-SPS 1900K Ox ide+glass-Lev ine et al.[5] 1900K depletion Lev ine et al.[5] 1873K modelox ide 1873K model-ox ide+glass 1873K model-depletion  Fig. 5.  (a) Model compared with oxidation weight gain measured time during oxidation of ZrB2-20 vol%SiC samples as a function of in static air at 1873 K, by Carney et al.10 and at 1900 and 1600 K et al.5 The  by Levine  solid lines are model predictions  for  the  two  temperatures.  (b) Model compared with (oxide + glass)  reported  oxide  thickness  (open),  total  thickness  (filled), and internal depletion  thickness  (dash)  as marked.  In  both  (a)  and  (b),  the  squares  represent Carney’s hot-pressed samples,  the diamonds Carney’s SPS  samples, and triangles data from Levine et al.  January 2012  Oxidation Modeling of SiC-Containing Refractory Diborides  345          \\x0c', '325 min.55  From their metallographic  images  along with  EDS maps,  the  thicknesses of  the oxide  layer,  glass  layer,  and depletion layer were obtained. These are plotted along  with model predictions in Fig. 8(a). A bilinear approximation the measured temperature-time profile, used as  of  input  for  the model,  is  shown as  inset. More  recently, Savino et al.  have  conducted a  similar study at higher jet.56 From the  surface  tempera tures using the arc  reported thermal history  and calculated scale  thicknesses  from the published micro graphs,  the  plot  shown  in Fig. 8(b) was  generated which  shows the model predictions compared to the data. SiC-HfB2 et al.12 In  Data  on  arc jet tested by Carney11  samples  of  have  been  reported  on  and Gasch  particular,  Carney  reported  on  scale  thicknesses  for  both  furnace exposed and arc-jet-exposed samples of  the  same batch of  sample under  the  same  temperature  (1773 K) and duration.  Figure 8(c)  shows  the model predictions compared with fur nace data and the arc-jet data. Figure 8(d)  shows  the data  from Gasch et al., who conducted two tests: one at  surface  temperature of 1803 K, and another at 1963 K. The model  predictions for both these temperatures are shown.  In general,  the model underpredicts the kinetics for experi mental  data  on  samples  exposed  to  arc  jet. The model  is  much closer at  lower  temperatures. The data from furnace exposed samples  show slower kinetics and are in reasonable  correspondence with the model.  IV.  Discussion  (1)  Model Strength and Weaknesses  A model has been presented to interpret  the oxidation kinet ics of SiC-containing refractory metal diborides. The model  includes  the  effect of viscous flow of  the outer glassy scale  and evaporative  losses of volatile  species under a boundary  layer  condition  established  by  ambient  fluid  flow velocity.  The model  predicts  the  total weight  of  oxygen  consumed,  sample weight  gain, weight  of  evaporated  species,  oxide  thickness,  glass  thickness,  depletion  layer  thickness,  and  recession of  the substrate. The model uses as  input parame ters,  the exposure temperature and time or a thermal profile,  ambient  oxygen  partial  pressure,  gas  chemistry,  and  fluid  (a)   (b)   1.E 03  1.E 02  1.E 01  1.E+00  4.0E 04  5.0E 04  6.0E 04  7.0E 04  W  e  i  g  h  t  g  a  i  n  ,  g k  /  m  2  1/ Temperature,K  2222  2000  1818  1666 K  Sample + Crucible  Sample  Wt. O2  Wt. Gain  Wt. Evaporated  1.00E 06  1.00E 05  1.00E 04  1.00E 03  4.00E 04  5.00E 04  6.00E 04  7.00E 04  S  c  a  l  e  T  h  i  c  k  n  e  s s  ,  m  1/ Temperature, K  2222  2000  1818  1666 K  Total Scale  Deple(cid:415)on  Oxide + deple(cid:415)on  Fig. 6. Comparision of model with experimental data on oxidation kinetics of HfB2-20 vol%SiC in the high-temperature regime (up to 2173 K) in static air showing weight gain in (a) and scale thicknesses from the work of Carney.11 The solid lines  in (b). All  the data are  are model predictions. sample + crucible  In  (a),  both  sample  weight  (open)  and  (filled)  weights  are  shown,  indicating  flow  is  significant  above  2000 K.  Similarly  in  (b),  total  scale  thickness  (filled)  and  oxide  scale  thickness  (open)  are  shown.  Due  to  experimental ambiguity in demarcation between depletion and oxide scale + depletion  scale,  oxide  thickness  are  shown  for  data  and  model.  (a)   0  0.05  0.1  0.15  0  100  200  300  400  W  e  i  g  h  t  g  a  i  n  ,  K  g  /  m  2  Time, min  1773K  (Hot Pressed)  Wt O2  Wt gain   (b)  0  0.05  0.1  0.15  0  100  200  300  400  W  e  i  g  h  t  g  a  i  n  ,  K  g  /  m  2  Time , min  1873K  (SPS processed)  Wt O2  Wt gain  (c)  0  0.05  0.1  0.15  0  100  200  300  400  W  e  i  g  h  t  g  a  i  n  ,  K  g  /  m  2  Time, min  1873K  (Hot Pressed)  Wt O2  Wt gain  Fig. 7.  Oxidation data on different batches of hot-pressed (a and c) and SPS-processed samples (b) of 20 vol%SiC-HfB2 composition, collected at 1773 and 1873 K by Sevener.20 The solid lines are model  predictions for  the weight gain and weight of oxygen consumed;  the  difference arises from evaporation of volatile oxides.  346  Journal of  the American Ceramic Society—Parthasarathy et al.  Vol. 95, No. 1  \\x0c', 'velocity. The material parameters  for  input are the size and  volume  fraction of SiC,  sample  length along the fluid flow  direction. The model calculates oxide + glass region, but glassy layer is assumed to be a fraction of  the activity of boria in the  the activity of boria at  the external  the inner  region.  This was  the only fit parameter. The model used a CO/CO2 counter diffusion mechanism for oxygen transport within the  depleted region.  In Fig. 2,  the model  is  seen to be  in agreement with the  kinetics  for oxidation of pure SiC in dry oxygen, although  there is  significant  scatter  in the data. This agreement  shows  that oxidation kinetics of SiC is limited by diffusion of molec ular oxygen across  the silica scale, assumed in the model.  It  further  lends credibility to the assumption that  the escape of  CO, one of the oxidation products,  is not rate limiting.  From Fig. 3,  the  effect of volume  fraction of SiC on the  kinetics of ZrB2-based UHTCs with the model for the two different  is  seen to be  in agreement  sources of experimental  data. The model predictions  for  the  total oxygen consumed  are  shown by dotted lines, and the predicted sample weight  gain by solid lines. The loss of material  from the surface due  to  evaporation  and  fluid  flow accounts  for  the  difference  between the  two. Figure 3(c)  confirms  that  the model  cap tures the parabolic time dependence quantitatively.  In Fig. 4(a),  the  experimental data on weight gain agree  well with the model, and the model predicts accelerated evap oration at  temperatures above 1873 K. In Fig. 4(b),  the mea sured scale  thicknesses are  in agreement with model within  the bounds of  scatter  in the  experimental data. The  scatter  arises  from the  variations in the flow of in Carney.10 The  externally  glassy et al.14  layer,  as  detailed  data  of Zhang  were obtained using a zirconia furnace, where the possibility  of contamination resulting in accelerated oxidation cannot be ruled out. The model predicts the depletion layer to be 2 lm 3 lm at  at  1673 K rising  to only  1873 K. Figure 5  shows  that the evolution in time of weight gain and scale thicknesses agree well with the model. The data of Carney10 at 1873 K and the data of Levine et al.5 at 1600 and 1900 K are used  to compare with the model. The  two symbols  for data of  Carney refer  to samples processed by SPS and hot pressing.  The data of Levine et al.  fall slightly below these. The deple tion depths  reported by Levine  et al. at 1900 K are higher  than the model prediction, while depletion was not observed  by Carney at 1873 K.  The model predicts a finite but small depleted zone. The variations. A 10-lm thick  experimental  reports  show large  depletion  zone was  reported by Fahrenholtz et al.14  at  1773 K in  30 min, whereas  Zhang  report  no  depleted  zone  below  1873 K.  Similarly,  no  clear  depleted  zone  was 1673-  observed  and  reported  by  Carney  after  2.5 h  at  1873 K. However, extensive EDS and FIB/TEM analyses did zone which was 2 lm at  reveal a partially oxidized internal 1673 K and 14 lm at 1873 K. As  the actual  fraction of SiC  oxidized in this partially oxidized zone was not measured,  it  is difficult  to compare with this model.  In summary,  some  ambiguity remains with respect  to internal depleted zone for mation in these materials. the HfB2-SiC data of Carney11 are rationalized well by the model up to 2073 K, but there is a significant  In Fig. 6,  increase  in measured oxidation compared to the predictions  at 2173 K. The  reasons  for  this  remain unknown, but per haps contamination from the stabilized zirconia heating ele ments may have  caused this discrepancy because  significant  Ca was detected in EDS of  the scales. The scales were found  (a)  (b)   (c)  (d)   0.0E+00  5.0E-06  1.0E-05  1.5E-05  2.0E-05  2.5E-05  3.0E-05  3.5E-05  4.0E-05  0  2  4  6  8  h T  i  k c  n  s s e  ,  m  Time, min  Montverde, Savino [33] Arc jet : ZrB2-15SiC   2193K, 325 m in  oxide  deple(cid:415)on  oxide + glass  0.0E+00  2.0E 05  4.0E 05  6.0E 05  8.0E 05  1.0E 04  1.2E 04  1.4E 04  0  2  4  6  8  10  T  h  i  c  k  n  e  s s  ,  m  Time, min  deple(cid:415)on  oxide  oxide + glass  Savino et al., [34] ZrB2 15SiC : Arc Jet 1573K 2053K, 540 min  oxide  deple(cid:415)on  oxide + glass  0.0E+00  5.0E 06  1.0E 05  1.5E 05  2.0E 05  2.5E 05  0  2  4  6  T  h  i  c  k  n  e  s s  ,  m  Time, min  Carney [11] HfB2 20SiC 1773 K, 4min  Furance  Arc Jet  Oxide + glass  Oxide  0 .0E+00  5.0E 06  1.0E 05  1.5E 05  2.0E 05  2.5E 05  3.0E 05  3.5E 05  4.0E 05  0  5  10  15  T  h  i  c  k  n  e  s s  ,  m  Time, min  1963K 10 min  1803K 10 min  1963K Total Scale  Oxide Scale  1803K Total Scale Oxide Scale  Gasch, Johnson [12] HfB2 20SiC : Arc Jet  Fig. 8.  Results  from arc-jet  testing done on ZrB2-SiC (a and b) and on HfB2-SiC (c and d) are shown along with model predictions reported thermal history used in the experiment (shown as inset in a and b) (data from ref. 11,12,55,56). In general, the model is  for  the  seen to  underpredict  the oxidation kinetics, with the discrepancy increasing with temperature of exposure. In one case, shown in (c), a furnace oxidation  and arc-jet  run were conducted for  the same temperature and time. The model  is  found to be much closer  to the furnace-based data and well  below the arc-jet data.  January 2012  Oxidation Modeling of SiC-Containing Refractory Diborides  347    \\x0c', 'to have  a  complex morphology  at  temperatures of  1973 K  and above. The  expected depletion layer was not observed; regions as detailed in Carney were noted.11  several different  Thus,  the plot only shows  the  combined thickness of oxide  and  depletion  layer,  and  the  total  scale  thickness  for  the  experimental data. Within experimental scatter,  in Fig. 7,  the  model captures the evolution in time of weight gain at 1773 and 1873 K, as reported by Sevener.20  Figure 8 reveals the key inadequacy of the model. The oxi dation kinetics during arc-jet testing were measured and reported by Montverde and Savino for ZrB2-15SiC under two different heat flux conditions.55,56 The actual thermal profiles  reported were used in the model. The model  fails  to capture  even the  correct order of magnitude of  scale  thicknesses at  high heat fluxes, while it is reasonable for low heat flux condi tions. The internal depletion is Similarly for HfB2-20SiC, ics by a significant margin. In one reported work, both furnace  considerably underpredicted.  the model underpredicts the kinet oxidation and arc-jet oxidation were  conducted at  the  same  temperature  and  time,  viz.  1773 K for  4 min. Figure 8(d)  shows that once again,  the model captures arc-jet data better  at lower temperatures, but not at higher temperatures using data from Gasch et al. on HfB2-20SiC.12 The model prediction is close for the furnace oxidation data, but far below the  arc-jet data. Clearly, aspects of oxidation kinetics  that  take  place during arc-jet conditions are not captured in the model.  (2)  Future Directions  The model  is  clearly  lacking  in capturing high-temperature  oxidation kinetics under  arc-jet  conditions. Several possible  mechanisms  can  be  suggested  that  require  further  analysis  and  experimentation. The  dissociation  of  gases  into mon atomic oxygen and nitrogen under arc-jet conditions  is well known.(for  example,  see  ref.  57) However,  the  oxidation  mechanisms under such conditions have only been studied by a few.58,59 These studies  show that oxidation is  indeed faster  under  dissociated  oxygen  conditions;  however,  the  reasons  for  this have not been established. The recombination of  the  monatomic gases at  the material  surface  can release  energy  that  is sufficient  to raise the material surface temperature.  In  the  furnace  oxidation  tests  using monatomic  oxygen,  the  material  surface  temperature  could not be measured. How ever,  during  arc-jet  testing,  the material  temperature was  measured, and the model  is unable to capture the oxidation  kinetics with  these measured  temperatures,  indicating  that  some other mechanism might play a role. The possibility of  enhanced oxidation kinetics  from permeation of oxygen in  monatomic  form has been suggested, but no direct  experi mental  evidence  exists. A modeling work by Li  et al. based  on first principles showed that  the dissolved concentration of  monatomic oxygen in borosilicate glasses can be higher than 1773 K.60  that  of  diatomic  oxygen,  at  temperatures  above  Given that  the  arc-jet  test data  are  closer  to the model  at  lower  temperatures  (below 1900 K) and deviate  significantly  at higher  temperatures,  the  enhanced  permeation  by mon atomic oxygen in borosilicate glasses at higher  temperatures  is  a possibility. Monatomic oxygen would also change  the  thermodynamics  of  oxidation  at  the ZrB2/oxide investigations that could con interface.  Theoretical and experimental  firm this hypothesis are suggested for  future work.  It  is also  suggested that  the possibility of active oxidation of SiC from  the surface by monatomic oxygen be investigated. The effect  of mechanical  shear at high velocities has been neglected in  this work;  however,  the model  predicts  evaporation  to  be  high enough to predict  the absence of  external glass under  arc-jet conditions.  Another  limitation of  the model  is  its  inability to predict  the  depletion  layer  thicknesses  in  all  the  reported works.  However,  there  is  considerable disagreement between differ ent works on the extent and even existence of  the depletion  layer under  identical conditions  for  similar or  identical com positions. The model shows clearly that counter diffusion via  the CO/CO2 mixture as oxygen transport medium is and can lead to a depleted region. However, the extent  tenable  is not  as high as  in some of  the reports, while it  is consistent with  others. Some uncontrolled variable, such as humidity or sam ple porosity, within the experiment might be responsible for  this. Future studies  that clarify this might be useful. Finally,  the effect of mechanical  forces that  induce shear of  the exter nal glass by viscous flow has been ignored in this work and  may be worth examining in future models.  V.  Summary  A mechanistic model has been built  to interpret experimental  data on the oxidation kinetics of SiC-containing diborides of regime of 1473-2473 K. The  Zr and Hf  in the  temperature  model  uses  available  viscosity,  thermodynamic  and  kinetic  data  on  borosilicate  glasses,  and  uses  a  logarithmic mean  approximation  for  compositional  variations.  The  internal  depletion region of SiC is modeled using counter diffusion in  CO/CO2 after furnace oxidation of pure SiC in oxygen, SiC-containZrB2 SiC-HfB2 UHTCs were found to agree reasonably well with the model.  as  oxygen  transport mechanism. Data  obtained  ing  with  varying  volume  fractions,  and  The model predictions  include weight of oxygen consumed,  sample weight  gain,  external  glassy  layer  thickness,  inner  oxide  scale  thickness,  recession,  and  depletion  layer  thick ness.  Some  depletion  layer  thicknesses  and  data  obtained  using  arc-jet  tests  fall well  outside  the model  predictions,  indicating  that  some  aspects  of  arc-jet  conditions  are  not  captured in this model.  Acknowledgments  We  acknowledge  useful  discussions with Dr.  I. Talmy  of NSWC, MD.  It  is  also a pleasure  to acknowledge  several useful  suggestions made by Prof.  E. Opila of Univ. of Virgina, during manuscript preparation.  References  1M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, “Oxidation-Based Materials Selection for 2000°C+ Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience,” J. Mater. Sci., 39, 5887-904 (2004). 2M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. 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},{
  "_id": 132,
  "PDF": "New Ablation-Resistant Material Candidate for Hypersonic Applications Synthesis, Composition, and Oxidation Resistance of HfIr3-Based Solid Solution..pdf",
  "Text": "['Cite This: ACS Appl. Mater.  Interfaces 2018, 10, 13062−13072  Research Article  www.acsami.org  New Ablation-Resistant Material Candidate for Hypersonic Applications: Synthesis, Composition, and Oxidation Resistance of HfIr3‑Based Solid Solution  Victor V. Lozanov,*,†  Natalya I. Baklanova,†  Natalia V. Bulina,†  and Anatoly T. Titov‡  †  ‡  Institute of Solid State Chemistry and Mechanochemistry SB RAS, Kutateladze Street 18, Novosibirsk 630128, Russian Federation  V. S. Sobolev Institute of Geology and Mineralogy SB RAS, Koptyug Avenue 3, Novosibirsk 630090, Russian Federation  *S  Supporting Information  ABSTRACT: The peculiarities of the solid-state interaction in the HfC−Ir system have been studied within the 1000−1600 °C temperature range using a set of modern analytical techniques. It was stated that the interaction of HfC with iridium becomes noticeable at temperatures as low as 1000−1100 °C and results in the formation of HfIr3-based substitutional solid solution. The homogeneity range of the HfIr3±x phase was evaluated −HfIr3.36. The durability of and refined as HfIr2.43 the HfIr3-based system under extreme environmental conditions was studied. It was shown that the HfIr3-based material displays excellent ablation resistance under extreme environmental conditions. The benefits of the new designed material result from its relative oxygen impermeability and special microstructure similar to superalloys. The results obtained in this work allow us to consider HfIr3 as a very promising candidate for extreme applications.  KEYWORDS:  iridium, hafnium carbide,  solid-state reaction,  intermetallics, ablation resistance  1.  INTRODUCTION  The search for new high-temperature materials has gained increasing impetus over the recent years. The environmental conditions under which the materials must operate become more and more severe. The trend toward the optimization of efficiency in modern gas turbines, combustors, and nose tips requires the development of high-strength materials with structural stability and corrosion resistance at temperatures than 2000 °C during operation.1,2 higher Several classes of materials have been proposed as the most appropriate materials for these purposes. First, these are monolithic ceramics based on hafnium-containing compounds such as HfC, HfN, and HfB2. The product of the oxidation of these compounds is hafnium dioxide, which is among the most refractory oxides and has the lowest vapor pressure at T = 2400 K (P = 1.5 × 10−7 atm3,4).5−7 Much attention was also paid to the research and development of the materials based on HfB2 or ZrB2 in conjunction with SiC. They were identified as the most appropriate oxidation-resistant refractory materials at temperatures above 1800 °C.1,2 Another class of promising materials for ultrahigh-temperature applications , namely , platinum group metal-based intermetallic compounds, especially those containing hafnium, have been proposed and comprehensively studied by YamabeMitarai and co-workers.8−13 It was shown that the platinum group metal-based superalloys, in particular, Hf−Ir-based superalloys, exhibit very promising mechanical properties because of their special microstructure composed of the  precipitates of the MeIr3 phase with the L12 structure embedded in the face-centered cubic (fcc)-compatible iridium matrix phase. A key component of these superalloys is the Engel−Brewer MeIr3 compound which is often called the phase. One of the promising intermetallics, HfIr3, has a set of other outstanding properties, such as high melting temperature (2470 °C), high bulk modulus value (270 GPa), and a hardness value of 6.05 GPa.14 In addition, HfIr3 may exhibit very good oxidation resistance at high temperatures because one of the components, hafnium, is oxidized with the formation of a refractory hafnium dioxide having the lowest vapor pressure at T = 2400 K, whereas the other component of intermetallics, iridium, forms gaseous iridium oxide with a very low regression velocity even at 2000 °C.15,16 Although the hafnium−iridium intermetallics are of great interest mainly as promising candidates for ultrahigh-temperapplications,17−19 ature there are other exciting areas of related materials.20−22 Here, applications of HfIr3 and one electrocatalysis.20−22 On could mention the basis of the literature data, one can emphasize that HfIr3 became more and more involved in different application areas. To extend the scope of application of this promising compound, a detailed study of HfIr3 and the development of synthetic methods must be undertaken. Earlier, it has been reported that HfIr3 can be  Received:  January 25, 2018  Accepted:  March 28, 2018 Published: March 28, 2018  © 2018 American Chemical Society  13062  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  Downloaded via FLORIDA INTL UNIV on January 8, 2020 at 20:45:43 (UTC).See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.\\x0c', 'ACS Applied Materials & Interfaces  formed through the solid-state reaction of hafnium carbide with iridium.17−19,23−26 Other approaches involve arc melting of hafnium and iridium ingots under an argon atmosphere8−14 and gas-phase deposition.18,27,28 To obtain a homogeneous product by arc melting, the procedure must be repeated five to six times. As to the formation of hafnium−iridium intermetallics by chemical vapor deposition (CVD), the attempt failed because instead of Hf−Ir of the formation of a ternary HfIrSi product intermetallics.28 The solid-state reaction of hafnium carbide with iridium remains one of the most convenient and simple methods to synthesize HfIr3. It was stated previously that HfIr3 is the only intermetallic phase formed from the reaction of iridium with hafnium carbide at approximately 2000 °C, independent of the HfC−Ir ratio.17,18,24,26 The authors explained this result by a Ir−HfC interface. The very low activity of hafnium at the results of the Ir−HfC diffusion couple experiments also showed that iridium diffusion is faster than the hafnium diffusion of HfIr3 compounds. Despite the continuously growing number of reports on synthesis, microstructure, characterization, and performance of Hf−Ir intermetallics, an elaborate study of the most important HfIr3 intermetallics has been lacking. The purpose of this work is to study the peculiarities of the solid-state interaction in the HfC−Ir system under different experimental conditions. The response of the HfC−Ir system in terms of phase composition, homogeneity region, microstructure and morphologies on temperature, exposure time, different HfC−Ir synthesis ratios, and others are discussed. An additional purpose of this work was to study the durability of the HfIr3-based system under extreme environmental conditions.  2. EXPERIMENTAL SECTION  2.1.  Initial Materials and Preparation of the Mixed Powders.  Hafnium carbide powders (TU 6-09-03-361-78) and iridium (GOST 12338-81; purity not less than 99.96%) were used as the initial substances. The as-prepared HfC powder was ground in an agate mortar for 5 min. The as-prepared (by the manufacturer) HfC powder and that preliminarily ground by us were packed separately in grafoil boxes (graphite tape, TU 5728-003-93978201-2008, carbon content not less than 99.5%), placed in a high-vacuum high-temperature furnace (SNVE-1.3.1/20; “PRIZMA”, Russia) and heat-treated to 1600 °C at a rate of 640 °C/h, kept at this temperature for 4 h, and then cooled at a rate of 300 °C/h. Iridium powder (ground in the agate mortar and as-prepared by the manufacturer) was heat-treated similar to the hafnium carbide powder. To study the mechanism of the interaction of hafnium carbide with iridium, several mixtures of the initial powders were prepared and heattreated under different conditions. The products formed in the HfC− Ir system were investigated depending on the treatment temperature, the HfC−Ir the exposure time at a given temperature, ratio, and the type of pretreatment (with or without grinding in the agate mortar). The experiments were carried out in the heat-treatment mode described as follows. The as-prepared 1:1 and 1:3 HfC−Ir mixtures were placed in grafoil boxes, heated up to a given temperature, T, under a vacuum of 10−5 Torr at a rate of 640 °C/h, and kept at a given temperature for different times, τ, from 1 to 4 h (Table 1).  Table 1. Experimental Details of Ir Powder Mixturesa  the Treatment of HfC and  aHere, τ is the exposure time at a given temperature.  Research Article  2.2. Sample Preparation for Ablation Resistance Tests and  the Testing Procedure. To study the ablation resistance of hafnium iridide-based systems, the as-prepared powders of HfC, iridium, and a small quantity of silicon powder (15 vol %) as the sintering additive were carefully mixed. Earlier, Johnson et al. stated that the addition of −SiC ceramics drastically changes iridium to HfB2 the microstructure and provides an improved oxidation resistance of ceramics.29 A powder mixture was loaded into a 30 mm diameter steel die and compacted as a pellet having a size of Ø 30 × 10 mm. Then the pellet was placed in a high-temperature high-vacuum furnace (SNVE-1.3.1/ 20; “PRIZMA”, Russia) and slowly heated up to 1600 °C, kept at this temperature for 2 h, and cooled at a rate of 300 °C/h. The ablation resistance of the samples was tested using a plasma range was 120−400 A; generator EDG-200M. The working current the voltage of the plasma generator was 120 ± 1 V. The samples were exposed to the arcjet in perpendicular geometry, and the maximum temperature of the ablation center reached up to 2200 °C, which was measured with an optical pyrometer focused on the hot zone of the specimen. The tests lasted for 1000 s. 2.3. Characterization. The X-ray diffraction (XRD) patterns of the powders were recorded with a D8 ADVANCE (Bruker AXS, in the θ−θ geometry using Cu Kα Germany) powder diffractometer (λ 1 = 1.54056 Å, λ 2 = 1.54439 Å) radiation, equipped with a onedimensional LynxEye detector and a Kβ filter. The XRD patterns were collected in the interval 30° < 2Θ < 130° with a step of Δ2Θ = 0.009° and a counting time of 177 s per step. Quantitative phase analysis and lattice parameter refinement were performed by the Rietveld method using the Topas 4.2 software (Bruker AXS, Germany). The instrumental contribution was calculated using the method of fundamental parameters.30 The information about the structure was taken from the Inorganic Crystal Structure Database (FIZ Karlsruhe, Germany, 1996). The observations of the morphology of the samples and the examination of the local elemental compositions were conducted with high-resolution scanning electron microscopes MIRA3 LMU (TESCAN, Czech Republic) and TM-1000 (Hitachi Ltd., Japan), coupled with energy-dispersive X-ray spectroscopy (EDXS) instruments INCA Energy 450 XMax 80 and Swift-TM (Oxford Instruments Analytical Ltd., GB), respectively. To observe the morphology of the powders, the samples were coated with a conducting chromium layer. To determine the elemental composition, the powders under investigation were packed with epoxy, cross-sectioned, and polished using 1 μm diamond suspension (Monosyn Duo, Synercon, Germany). The variation coefficient characterizing the reproducibility of a single determination is found to be 1% for EDXS within the compositional range of the main components (content more than 10%). The Fourier transform Raman spectra of the products were recorded using a Bruker RFS 100/S spectrometer equipped with a Nd:YAG laser operating at an excitation wavelength of 1064 nm. The laser output was 100 mW. For each spectrum, 1000 scans were accumulated. This was sufficient to obtain the spectra with a low noise/signal ratio required for the fitting procedure. The samples were prepared as mixtures with KBr at a 1:50 ratio. The deconvolution of the spectra was performed with Fityk 0.9.8 software.31 Voigt function profiles were used for the peak-fitting procedure.32,33 The HfC−Ir pellets before and after the ablation tests were analyzed using XRD and scanning electron microscopy (SEM)/EDXS. The procedures analogous to those described above to study the powders were used to investigate the morphology, phase, and elemental composition of the pellets.  ■ RESULTS AND DISCUSSION  Hafnium Carbide and Iridium Powders: The Effect of  Grinding and Temperature. Before studying the solid-state reaction of HfC and Ir, it was very important to clarify what features are related to the solid-state reaction immediately and what features are related to the effect of temperature on the initial substances. The effect of both grinding and temperature on the microstructure of the as-prepared HfC powder was  13063  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  \\x0c', 'ACS Applied Materials & Interfaces  studied using X-ray and SEM analyses. According to the XRD data, no noticeable changes occur in the XRD pattern of HfC after grinding in the agate mortar. The thermal treatment at 1600 °C for 4 h leads to hardly noticeable narrowing of the Xray peaks of the as-prepared and ground HfC powders (Supporting information, Figure S1), which corresponds to an increase in the crystallite size from 140 to 280 nm with temperature, as determined with the quantitative X-ray analysis. The effect of grinding on the microstructure of iridium powder is more significant, compared with that of HfC powder, which can be clearly seen from the X-ray data (Supporting Information, Figure S2) and SEM observations (Figure 1).  Figure 1. SEM images of iridium powder: (a) as-prepared by manufacturer and (b) ground with subsequent thermal treatment.  the  According to the XRD data, the X-ray peaks in the range of 2Θ = 120−130° are broadened, compared with those for the asprepared iridium powder (Supporting Information, Figure S2). This result can be associated with a decrease in the grain size and an increase in the number of surface defects after grinding (Supporting Information, Figure S3a,b).34−37 At the same time, no visible traces of plastic deformations (such as steps on the grain surface) have been observed (Supporting Information, Figure S3b). The thermal treatment at 1600 °C of the as-prepared iridium powder resulted in the narrowing of X-ray peaks. One can note temperature as high as 1600 °C is 500 °C that the treatment above the Tammann temperature for iridium. Hence, the observable slight narrowing of X-ray peaks can be assigned to the recrystallization of iridium (Supporting Information, Figure S3). The thermal treatment of the preliminarily ground iridium powder results in strong distortions of the X-ray patterns, which are especially noticeable in the range of 2Θ = 120−130°. The distortions are associated with the formation and accumulation of the macro stresses, which are also confirmed by the SEM images. Indeed, the appearance of numerous steps on the surface of the grains is detected (Figure 1b).  3.2. Hafnium Carbide and Iridium Powder Mixture: The Effect of the Temperature. The mixtures with HfC−Ir  = 1:1 and 1:3 ratios were studied in detail. The XRD patterns of the HfC−Ir mixture with the 1:1 ratio depending on the temperature are presented in Figure 2. The analysis of the XRD patterns detected that several phases are present, namely, HfIr3 intermetallics, HfC, and iridium. No intermetallic compounds other than HfIr3 are formed, despite the fact that there is a set hafnium−iridium intermetallics the Hf−Ir system.38,39 of in The first features of the appearance of the HfIr3 phase become noticeable at 1100 (Figure 2a). According to the quantitative X-ray analysis, the yield of HfIr3 at this temperature is about 3% (wt). One can note that the yield of the target phase, measured with the quantitative XRD analysis, appears to be overestimated because the major part of iridium is undetectable because of the absorption of X-ray radiation.  °C  Research Article  This suggestion was supported by the calculation of material balance for iridium. As the quantity of unreacted iridium decreases, the uncertainty of the determination of iridium by the quantitative XRD analysis drops too. One can consider that the large values of intermetallic phase yield are close to true values. With temperature, the intermetallic phase yield uniformly increases and reaches 34% (wt) at 1600 °C (Figure 3). As one can see from Figure 3, the reproducibility of the obtained values is not so high, and there is a noticeable deviation in the yield values. Usually, such a poor reproducibility of solid-state reactions can be connected with the scarce reproducibility of the state of solid surfaces and the degree of contact between them.40 It is noteworthy that the most significant changes were found to occur for those X-ray peaks that are associated with the HfIr3 phase. The XRD peak for 2Θ = 116−118° has a complicated shape for the temperature of 1200 °C and higher. Thus, in the range of 2Θ = 116−118°, an appearance of the (331) peak belonging to the HfIr3 phase in the XRD pattern for 1200 °C and a weak splitting of the peak into two broad peaks are observed. Starting from 1300 °C, the splitting becomes more obvious, and the distance between the maxima of the peaks increases with temperature (Figure 2b). The most plausible reason for the observable features is the formation of HfIr3based continuous solid solutions. The lattice parameter of the HfIr3±x phase, a, was calculated from the experimental XRD patterns of the products obtained at different temperatures. The parameter a, calculated from the XRD pattern, associated with the maximum at 2Θ ≈ 116.5° (to the left of the value for stoichiometric HfI3), was proved to be equal to 3.954 Å for the product obtained at 1200 °C. The parameter decreases to 3.948 Å for the product obtained at 1600 °C. It is noteworthy that the observable peculiarities of the XRD patterns in the large-angle range for HfIr3±x are reproducible. As to the other phases present in the XRD pattern, namely, unreacted Ir and HfC phases, one can note that the position of the Kα1 and Kα2 doublet for the (222) peak belonging to the HfC phase is not changed with temperature (Figure 2, inset c). This means that iridium does not dissolve in hafnium carbide within the whole temperature range. On the contrary, the positions of the XRD peaks of iridium are shifted toward the small-angle direction, which can be especially clearly observed in the large-angle range of the pattern. The shift is accompanied by a decrease in peak intensities (Figure 2, inset c). This suggests that the lattice parameter of iridium is increased, which can be related to the partial substitution of iridium atoms by hafnium atoms in the crystal structure of iridium. Because the radius of hafnium atom is larger than that of iridium (2.08 Å vs 1.80 Å), the substitution leads to an increase in the lattice parameter of iridium. In this work, the lattice parameter of the Ir(Hf) solid solution was determined to be equal to 3.847 Å (for 1600 °C). One can note that the observable peculiarities of the XRD the HfC−Ir = 1:1 mixture patterns for heated to different temperatures are well-reproducible. The role of the carbon phase eliminated in the course of the solid-state reaction of HfC with iridium will be discussed below. The XRD results for the HfC−Ir = 1:3 mixture showed that the regularities determining the effect of temperature on the phase composition of the products are the same as those the HfC−Ir deduced for = 1:1 mixture described above (Supporting Information, Figure S4). With temperature, the  13064  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  \\x0c', 'ACS Applied Materials & Interfaces  Research Article  the HfC−Ir = 1:1 mixture depending on the temperature and selected regions of Figure 2. Survey XRD patterns of the XRD patterns: inset (a) appearance of the strongest 111 peak of HfIr3 and inset (b) splitting of the (331) peak of the HfIr3±x phase because of the formation of the solid solutions. The dashed line marks the peak position of stoichiometric HfIr3 after Copeland;38 inset (c) selected region of the XRD pattern showing the unchangeable position of the Kα1 and Kα2 doublet for the (222) peak belonging to the HfC phase with temperature and a shift of the Kα1 and Kα2 doublet for the (220) peak belonging to the Ir phase with temperature.  yield gradually increases up to 41% (wt) at 1600 °C (Figure 3). The same phenomena, namely, (i) splitting of the (331) peak of the HfIr3±x phase into several peaks and (ii) a shift of the Kα1 and Kα2 doublet for the (220) peak belonging to the Ir phase with temperature, are observed. The lattice parameters of the HfIr3±x phase vary rather significantly. No intermetallic phases other than HfIr3±x were detected for the HfC−Ir = 1:3 mixture. The boundaries of the lattice parameter, a, determined from the XRD data of the HfIr3±x phase for both mixtures, are presented in Figure 4. The general trend is that the difference in a parameter values for each mixture increases with temperature. This means that the homogeneity range for the HfIr3±x phase broadens with temperature. Also, one can note that the lower curve corresponding to the upper boundary of the homogeneity of HfIr3+x falls down faster than the lower  13065  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  Figure 3. “Yield” of solid-solution HfIr3±x with the dependence on the temperature.  \\x0c', 'ACS Applied Materials & Interfaces  Figure 4. Boundaries of the lattice parameter, a, determined from the XRD data of the HfIr3±x phase for the HfC−Ir = 1:1 and 1:3 mixtures. The exposure time at a given temperature is 1 h.  boundary of homogeneity. This may be connected with a faster iridium diffusion compared with hafnium in HfIr3±x. For both the HfC−Ir = 1:1 and HfC−Ir = 1:3 mixtures, the shift of the XRD peaks of iridium and a decrease in peak intensities (Figure 2, inset c; Supporting Information, Figure S4, inset c) are observed, which can be an argument in favor of an increase in the lattice parameter of iridium because of the formation of the fcc-compatible iridium-based solid solution. One can propose that a hafnium atom having a larger atomic radius substitutes an iridium atom in the lattice. The largest lattice parameter of a substitutional Ir−Hf solid solution (3.849 Å) was observed in this work for the sample obtained at 1600  Research Article  °C. Although the HfC−Ir system was far from equilibrium, it can be proposed that the Hf(Ir)C solid solution is absent in this system, which is in good agreement with the phase equilibrium diagram for the Hf−Ir−C system proposed by Holleck.26  3.3. Hafnium Carbide and Iridium Powder Mixture: The Effect of Exposure Time at 1600 °C. A comparative  analysis of the XRD patterns obtained for the HfC and Ir mixture with the ratios of 1:1 and 1:3 exposed at 1600 °C in a vacuum of 10−5 Torr for different times showed that, first, no additional intermetallic phases other than HfIr3±x were observed (Figure 5 and Supporting Information, Figure S5). increases up to 72% (wt) for With time, the yield of HfIr3±x 1600 °C, τ = 4 h, and HfC−Ir = 1:3 ratio. Further, the X-ray peaks belonging to the HfIr3+x solid solution are more intensive, compared with those belonging to the HfIr3−x solid solution (Figure 5, inset a). Starting from 2 h exposure, the width of the homogeneity range is practically constant for both the 1:3 and 1:1 mixtures (Figure 6). The lattice parameters corresponding to the lower and upper boundaries of homogeneity for the HfIr3±x phase were determined to be equal to 3.947 and 3.928 Å and 3.945 and 3.926 Å for the 1:1 and 1:3 mixtures, respectively. One can note that the preliminary grinding of a mixture of HfC and Ir also promotes the significant yield of the HfIr3±x product because of the formation of macro defects in iridium, a decrease in grain sizes, an increase in surface energy and, as a consequence, the enhancement in the reactivity of iridium (Supporting Information).  Figure 5. Survey XRD patterns of the products obtained from the HfC−Ir = 1:3 mixture at 1600 °C depending on the exposure time and selected regions of the XRD patterns: inset (a) splitting of the (331) peak of the HfIr3±x phase because of the formation of the solid solutions. The dashed line shows the peak position for the stoichiometric HfIr3;38 inset (b) selected region of the XRD pattern showing the position of the Kα1 and Kα2 doublet of the (222) peak belonging to the HfC phase with time and the position of the Kα1 and Kα2 doublet for the (220) peak belonging to the Ir phase with time. Arrows mark an additional Kα1 and Kα2 doublet of the (220) peak of iridium, the lattice parameter being smaller than that of pure iridium.  13066  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  \\x0c', 'ACS Applied Materials & Interfaces  Research Article  3.4. Reaction of  Iridium with Hafnium Carbide. The  temperatures lower than the melting points of noble metals. Unfortunately, no explanations for this interesting phenomenon were proposed. In any case, this result needs a thorough investigation in future.  opportunity for the formation of the HfIr3 intermetallic phase with a high yield via a simple and convenient way using the solid-state reaction 1 between hafnium carbide and iridium was clearly demonstrated in this study. However, in the course of the study, some peculiarities of the solid-state interaction of hafnium carbide and iridium were revealed; therefore, a detailed discussion is necessary.  Figure 6. Lattice parameters homogeneity of the HfIr3±x mixtures.  for the upper and lower boundaries of phase obtained for the 1:1 and 1:3  According to the data reported by Eremenko et al.,39 the homogeneity range of the HfIr3±x phase was determined to be to 73−79% (at.) equal Ir. The authors mentioned that they calculated the homogeneity boundaries relying on the data of local X-ray spectroscopy, but unfortunately, no experimental lines of evidence were presented in their work. Earlier, Copeland et al.38 proposed that the homogeneity range can extend from stoichiometric to iridium-rich compositions beyond the stoichiometric composition (76.37 wt or 75 at. % Ir) at least to 78 (wt %). On the basis of the data on the lattice parameters obtained in this work, it is possible to estimate the phase41 composit ion range of the HfIr3±x (Supporting Information, calculations of solid-solution compositions). According to our calculations, the composition range could be 70.8−77.1% (at.), which corresponds to the composition −HfIr3.36. On range HfIr2.43 the whole, the width of the homogeneity range is in good agreement with that reported by Eremenko et al.,39 but the whole homogeneity area of the HfIr3±x phase determined in this work is slightly shifted toward a lower iridium content. It follows from Figure 5b that the peak positions for the HfC phase are unchanged for any exposure time at 1600 °C. As to the X-ray peaks belonging to iridium, their positions are also practically conserved and the intensities are diminished, which is the evidence of iridium consumption for the formation of the HfIr3±x phase. The lattice parameters of the Ir(Hf) solid solution are about 3.847−3.849 Å for this time interval. A comparative analysis of the XRD patterns in the range of 2Θ = 69−70° for the HfC−Ir mixtures exposed at 1600 °C for different times detected the appearance of additional smallintensity peaks for long exposure (Figure 5b). According to the quantitative XRD analysis, the lattice parameter of the iridium phase is 3.823 Å, which is smaller than that for pure iridium. The result could be explained by the substitution of an iridium atom (radius 1.80 Å) by any atom of a smaller size. The only component of the reaction system under investigation is the carbon atom (radius 0.67 Å). Therefore, it could be proposed with discretion that a decrease in the lattice parameter of iridium compared with pure iridium by 0.44% could be connected with the dissolution of carbon in iridium with the formation of a metastable solid solution. The observable result is not a single example of a decrease in the lattice parameter of a precious metal in the reaction with high-melting carbide. Earlier, Raub et al.42 also detected an obvious decrease in the lattice parameters of palladium and platinum in the course of the reaction of noble metals with tungsten carbide at  HfC  (s)  +  3Ir  (s)  =  HfIr  3(s)  +  C  (s)  (1)  It has been already shown in this work that the first features of the appearance of the HfIr3±x phase in the product were detected by the XRD analysis already at temperatures as low as °C. This 1100 is the lowest temperature among the values reported in the literature for the reaction of hafnium carbide with iridium. Previously, Strife et al.24 studied the interaction of hafnium carbide with iridium within a CVD-derived double coating and detected that HfC and Ir rapidly reacted with the in the temperature range of 1923−2400 K. formation of HfIr3 Kwon18 studied the reaction of an iridium foil with HfC in the same temperature range. The mentioned temperature range (1923−2400 K) is close to the melting point of iridium (2716 K), and therefore, it would be reasonable to propose that the rapid reaction between carbide and iridium was related to very high iridium mobility as the temperature approaches the melting point of iridium. The fact that hafnium carbide starts to decompose by iridium at 1100 °C is an evidence of its extreme instability in the presence of iridium even at low temperatures that are rather far from the me l t ing po in t o f ir id ium . On the bas is o f thermodynamic calculations, it was identified that the reaction between iridium and hafnium carbide can occur at room temperature.17,24 However, the diffusion of heavy atoms is rather slow, and therefore, kinetic limitations hinder the reaction at low temperatures. Partially, the limitations are taken off when the temperature approaches the Tammann iridium (T ≈ 1110 °C).43 The studies of temperature for the reactions of other precious metals, such as platinum and palladium, with the carbides of transition metals (W, Ta, Nb) at the temperature in the vicinity of the melting points of precious metals also confirm the fact that refractory carbides can be readily decomposed by precious metals.42 Only HfIr3±x intermetallics were obtained by the reaction of HfC with iridium, despite the existence of other intermetallics in the Hf−Ir system.17−19,24 This result can be explained by a the Ir−HfC interface. very low activity of hafnium at It is the iridium diffusion which dictates the growth of the HfIr3±x phase. Although both the initial substance (HfC) and the reaction product (HfIr3±x) have fcc unit cells, there is a rather large difference between the volume of HfC unit cell per atom (12.325 Å) and the volume of stoichiometric HfIr3 unit cell per atom (15.225 Å), that is, the solid-state reaction occurs with a significant increase in the unit cell volume. Hence, one can propose that the solid-state interaction between HfC and Ir appears to occur via the destruction of the HfC crystal lattice, followed by the formation of the HfIr3±x crystal lattice. Another question to be discussed is the formation of HfIr3±x solid solutions. As noted above, there are ambiguous data about the composition range of the HfIr3±x phase. Copeland38  13067  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  \\x0c', 'ACS Applied Materials & Interfaces  considered HfIr3±x as a phase with a narrow homogeneity range, whereas Eremenko et al.39 found that the homogeneity range spreads from 73 to 79 at. %. The results of this work based on the quantitative XRD analysis clearly demonstrate that, first, the HfIr3±x phase has a homogeneity range, and second, the homogeneity range is rather wide and spreads from HfIr2.43 to HfIr3.36. In addition, the presence of the HfIr3±x phase with different compositions can be observed by means of SEM/EDXS analysis (Figure 7).  Research Article  Figure 7. SEM images of the products obtained by the solid-state reaction of HfC with iridium. On the SEM images (a,b), the HfIr3±x phase with different compositions is presented: (a) selected areas with HfIr3−x composition and (b) selected areas with HfIr3+x composition.  The obtained results discover a new strategy for the synthesis of very stable, high-strength, and hard compounds such as HfIr3±x and other related compounds. Indeed, to synthesize HfIr3 intermetallics with a high yield, it is sufficient to mix hafnium carbide and iridium powders and to heat them at moderate temperatures in an inert atmosphere. One can note that this approach leads to the formation of an HfIr3±x product which is highly adherent to HfC (Figure 7a,b). It should be emphasized that the detection and evaluation of the phase composition of HfIr3 are of great importance. Indeed, this compound is still not sufficiently studied, and its properties are almost unknown. As a rule, there is a distinct dependence of properties on phase stoichiometry. This should be kept in mind when studying any property of the HfIr3±x phase. Let us consider the behavior of the iridium phase. As was mentioned above, there is a predominant diffusion of iridium atoms through the layer of the product, HfIr3±x, toward HfC. The hafnium atoms released from the crystal lattice move toward iridium. They are able to substitute the iridium atoms because the atomic radii of Ir and Hf are comparable (RIr = 1.80 Å vs RHf = 2.08 Å). In this case, the parameter of the iridium crystal lattice would be larger. Indeed, as was determined by the XRD analysis, the lattice parameter of Ir increases to 3.847 Å (for 1600 °C) that can be associated with the formation of Ir(Hf)-substitutional solid solution (Figures 2c and 5b). Special attention must be paid to the formation of the carbon phase in the course of the solid-state reaction of HfC with iridium. The SEM images of the intermetallic phase are presented in Figure 8a,b. On the whole, the grains of the product are poorly crystallized and have a rounded shape; however, some grains have well-crystallized habitus (Figure 8b). Also, one can see that transparent thin sheets are present on the surface of the product grains. As was proposed, these sheets could be related to the carbon phase. Indeed, intense Raman 1200−1600 cm−1 peaks in the region of and high-intensity overtones in the region of 2500−3000 cm−1 are detected in the  Figure 8. SEM images of the product obtained at 1600 °C for 1 h for HfC−Ir = 1:3: (a) thin carbon sheets on the surface of the grains; (b) well-crystallized grain of the product; and (c) Raman spectrum of the product (λ = 1064 nm) fitted by the Voigt function.  Raman spectra of the products derived under different experimental conditions (Figure 8c). The computer deconvolution of a Raman spectrum results in several constituents, two of which are centered at about 1280 (D-peak) and 1592 (Gpeak) cm−1, which can be considered as the Raman fingerprints of the carbon phase (Figure 8c). Carbon eliminated after the interaction of HfC with Ir appears to be a graphite-like carbon phase, although the formation of the multilayer graphene phase must not be excluded (Supporting Information).44,45 The presence of carbon as a free phase confirms the main conclusions deduced from the XRD analysis that no ternary phases in the Hf−C−Ir system, as well as a solid solution of carbon in HfC, are formed. On the other hand, the detection of iridium with the lattice parameter less than that for pure metal after the interaction of HfC with Ir could be a piece of evidence in favor of the formation of a metastable Ir-based solid solution, in which iridium atoms are substituted by atoms with smaller radius, in particular, carbon eliminated in the above-mentioned reaction (Supporting Information).  Resistance  of  3.5.  Ablation  HfIr3-Based Materials.  Figure 9 shows the cross-sectional microstructure at the center of the as-received pellet. One can see that this is a rather dense, sintered microstructure. It can be proposed that the addition of the HfC−Ir mixture a small amount of silicon powder in (Experimental Section) promoted the formation of low-melting Ir−Si eutectics, which served as a transient liquid-phase sintering agent.46 One can also see that several different phases are present. One of them is a continuous phase, which is composed of the IrSi phase according to the EDXS data. The other phases are discrete ones. According to the elemental analysis data, the discrete phases are HfC and HfIr3±x, as well as HfO2 (as the minor phase). The appearance of the last phase can be explained by the interaction of the hafnium-containing components with oxygen, which presents as an impurity in silicon. Examination of another interesting  the cross-sectional morphology detected feature, namely, a well-crystallized phase  13068  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  \\x0c', 'ACS Applied Materials & Interfaces  Research Article  the light-gray phase can be assigned to HfIr3±x, whereas the dark-gray discrete phase presents mainly iridium silicide as the main component. It is noteworthy that the microstructure described above spreads through the whole thickness of the pellet, and no features of gradient microstructure were observed. The reasons for the formation of the observable ordered microstructure are not yet clarified. Here, it should be noted that such a microstructure could form when the pellet was rapidly heated to ultrahigh temperatures and then cooled rapidly after the ablation test. The observable microstructure is analogous to that reported by Yamabe-Mitarai et al. for the solidified eutectic Hf−Ir−Si superalloys.10,11,46 As one can see from Figure 10c, the oxidized layer consists of monoclinic hafnium dioxide as the main phase. This suggests that the oxidation of hafnium iridide could occur in accordance with reaction 2.  HfIr  3(s)  2Ir  (s)  +  +  O  2(g)  =  HfO  2(s)  +  3Ir  (s)  3O  2(g)  =  2IrO  3(g)  ↑  (2)  (3)  It is noteworthy that the observable microstructure of the oxidized layer is in good agreement with the processes described by reactions 2 and 3. Indeed, iridium escapes during oxidation in the form of gaseous iridium oxide. Because iridium was disposed between the rows directed parallel to each other and composed of the grains of the dark-gray phase (IrSi), the evaporation of iridium led to the formation of empty areas between the rows and to a breakdown of the integrity of the oxidized layer. Nevertheless, a parallelism of oxidized rows is retained. One can propose that too high temperature of testing (2200 °C) promotes sintering of hafnium dioxide grains and keeping the grains together within the same layer. The obtained results showed that the HfIr3-based material displays excellent ablation resistance under oxidizing environments up to 2200 °C. The obvious benefits of the new designed material relative to the existing common UHTCs, for example, −SiC plus additives1,2) or silica/silicate formers (HfB2 the soca l led entropy-stab i l ized u ltrah igh-temperature mater ia ls (equiatomic multicomponent refractory carbides, nitrides, and carbonitrides of group IV, V, and VI transition metals47,48) result from the relative oxygen impermeability and special microstructure similar to superalloys.  4. CONCLUSIONS  In this work, a relatively simple approach was proposed to synthesize a high-melting compound HfIr3. This approach involves mixing of HfC and Ir powders, followed by heating to 1600 °C in an inert atmosphere. The peculiarities of the solidthe HfC−Ir state interaction in system have been studied within the 1000−1600 °C temperature range using a set of modern analytical techniques. It was stated that the interaction of HfC with iridium becomes noticeable at temperatures as low as 1000−1100 °C and results in the formation of HfIr3-based substitutional solid solution. Thus, HfC reveals instability to iridium in a wide temperature range. No intermetallic phases than HfIr3±x were detected for both the HfC−Ir = 1:3 other and 1:1 mixtures. The homogeneity range of the HfIr3±x phase −HfIr3 .36 . was eva luated and refined as HfIr2 .43 As the temperature and exposure time increase, the yield of the target phase, HfIr3±x, increases to 94% (wt). Thus, such highmelting intermetallics as HfIr3 can be synthesized in a high yield with the help of a very simple method, namely, simple mixing of HfC and iridium powders.  13069  ACS Appl. Mater.  DOI: 10.1021/acsami.8b01418 Interfaces 2018, 10, 13062−13072  Figure 9. Microstructure and phases of the cross section of the asreceived pellet: (a) cross-sectional morphology (backscattered electron image), (b) surface morphology (HfC cubes can be observed), and (c) compositional analysis of the selected areas of the central part of the pellet.  shaped as cubes. According to the EDXS data, this phase was identified as HfC. Comparison of the observed morphology of this phase in the pellet and that of the initial powder which was used for the preparation of the mixture clearly shows that the presence of cubes was not detected in the initial HfC powder. Therefore, we believe that the appearance of the welldissolution− crystallized HfC phase is connected with the crystallization of hafnium carbide during the preparation of the from the HfC−Ir(Si) mixtures. pellet The morphology of the cross section of the pellet oxidized at 2200 °C for 1000 s is presented in Figure 10. Two areas with distinctive morphology can be detected. The part exposed to arcjet is presented by strictly parallel rows composed of wellsintered grains. The thickness of the oxidized layer is approximately 100 μm. The remaining part of the pellet appears to be intact during ablation testing. The morphology of this part is presented by a well-ordered structure. The observable microstructure is composed of rows of dark-gray discrete phase embedded into a light-gray phase. There appears to be a very strong bonding at the “dark-gray−light-gray phase” interface because no features of detachment, cracks, pores, and so forth were detected. According to the EDXS analysis data,  \\x0c', 'ACS Applied Materials & Interfaces  Research Article  Figure 10. Morphology and composition of the cross section of the pellet oxidized at 2200 °C for 1000 s: (a) survey image (backscattered electron image); (b) enlarged image of the oxidized layer; and (c) XRD pattern of the oxide scale. The peak marked by an asterisk belongs to the strongest peak of the cubic HfO2; (d) surface temperature at the ablation center as a function of time during the test.  The durability of the HfIr3-based system under extreme environmental conditions was studied. It was proven that the HfIr3-based material displays excellent ablation resistance under extreme environmental conditions. The benefits of the new designed material result from its relative oxygen impermeability and special microstructure similar to superalloys, which provide the durability of the HfIr3-based material under extreme environmental conditions.  ■ ASSOCIATED CONTENT  *S  Supporting Information  The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b01418.  XRD patterns of HfC and Ir powders (the effect of grinding and temperature); SEM images of the Ir powder (after grinding and as-prepared with subsequent thermal treatment); XRD patterns of the mixture with HfC−Ir = 1:3 ratio (the effect of the temperature); XRD patterns of the mixture with HfC−Ir = 1:1 ratio (the effect of at 1600 °C); exposure time the effect of grinding on HfC−Ir mixtures; calculations of solid-solution compositions; assignment of Raman shifts; and general scheme of reactions of Ir and HfC (PDF)  ■ AUTHOR INFORMATION  Corresponding Author  *E-mail: lozanov.25@gmail.com. Phone: +7-(383)-233-24-10 ext. 1132. Fax: +7-(383)-332-28-47.  ORCID  Victor V. Lozanov: 0000-0003-1169-8224  Natalya I. Baklanova: 0000-0001-9930-9857  Natalia V. Bulina: 0000-0003-4751-0705 Anatoly T. Titov: 0000-0003-1320-4822  Notes  The authors declare no competing financial  interest.  ■ ACKNOWLEDGMENTS  The authors would like to thank PhD A. Utkin (ISSCM SB RAS) for useful discussion of the results of this work and D.Sci. I. Prosanov (ISSCM SB RAS) for help with the Raman study. This work was partially supported by the Russian Foundation for Basic Research (RFBR) under project mol_a 16-33-00613. This research was carried out within the State Assignment to ISSCM SB RAS (project 0301-2017-0001) and the State Assignment to IGM SB RAS (project 0330-2016-0013).  ■ REFERENCES  (1) Opeka, M. M.; Talmy, I. G.; Zaykoski, J. A. Oxidation-based selection for 2000°C+ hypersonic aerosurfaces: Theoretical materials considerations and historical experience. J. Mater. Sci. 2004, 39, 5887− 5904. (2) Ultra-High Temperature Ceramics: Materials for Extreme Environment Applications; Fahrenholtz, W. G., Wuchina, E. J., Lee, W. E., Zhou, Y., Eds.; Wiley: Hoboken, NJ, 2014. (3) Kazenas, E. K.; Tsvetkov, J. V. The Evaporation of Oxides; Nauka: Moscow, 1997 (4) Lozanov, V. V.; Baklanova, N. I.; Shayapov, V. R.; Berezin, A. S. 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},{
  "_id": 133,
  "PDF": "New insight into the formation and oxygen barrier mechanism of carbonaceous oxide interlayer in a multicomponent carbide.pdf",
  "Text": "['Received: 13 December 2019   DOI: 10.1111/jace.17143    |   Revised: 16 March 2020   |   Accepted: 26 March 2020  O R I G I N A L A R T I C L E  New insight into the formation and oxygen barrier mechanism of  carbonaceous oxide interlayer in a multicomponent carbide  Ziming\\xa0Ye1  Yalei\\xa0Wang1   |   Yi\\xa0Zeng1 |   Wei\\xa0Sun1   |   Xiang\\xa0Xiong1 |   Zhaoke\\xa0Chen1   |   Tianxiao\\xa0Qian1  |   Huilin\\xa0Lun1  |   Lijun\\xa0Zhang1  |   Ping\\xa0Xiao2  |     1State Key Laboratory of Powder  Metallurgy, Central South University,  Changsha, China 2School of Materials, The University of  Manchester, Manchester, UK  Correspondence  Yi Zeng, State Key Laboratory of Powder  Metallurgy, Central South University,  Changsha 410083, China. Email: zengyi001@csu.edu.cn  Funding information  National Natural Science Foundation of  China, Grant/Award Number: 51602349  and 51602351; Fundamental Research  Funds for the Central Universities;  Key research and development plan of  Hunan province, Grant/Award Number:  2018GK2061; Innovation-drive Project of  Central South University  ABSTRACT  Early transition metal carbides are considered to be superior candidate materials for  oxidizing environments at temperatures exceeding 2000°C. Generally, the remarkable oxidation resistance is largely attributed to a carbonaceous oxide interlayer (eg,  Hf-O-C, Zr-O-C, and Ta-O-C), located at the interface between the external oxide  layer and internal carbide (eg, HfC, ZrC, and TaC), acting as the primary oxygen barrier. However, the oxygen barrier mechanism of the carbonaceous oxide interlayer  remains unclear. Herein, through studying the oxidation behavior of a novel multicomponent carbide Hf0.5Zr0.3Ti0.2C in oxidizing environments up to 2500°C, the oxygen barrier mechanism of the carbonaceous oxide was recently revealed. We found  that the oxygen barrier resulted from the slow oxygen diffusion through the inner  grains of Hf-Zr-Ti-O due to the presence of carbon formed at the grain boundaries  because of the existence of compact external oxide layer, beneath which the Hf-Zr- Ti-O-C interlayer possesses much lower oxygen activity and temperature that allow  carbon to exist stably. This as-formed carbon strongly retarded the fast diffusion of  oxygen along the grain boundaries of oxides. Additionally, desirable synergisms of  the designed multicomponent system, particularly, the outward short-circuit diffusion of Ti, lead to the self-healing of the external oxide layer, evidently enhancing  integral protection performance against oxidizing environments.  K E Y W O R D S  carbonaceous oxide, interface, multicomponent carbide, oxygen barrier  1   |   I N T RO D U C T I O N  Early transition metal carbides (TMCs), possessing simple  metallic structures in which small carbon atoms occupy the  interstitial voids of densely packed metallic lattice, merge the  advantages of covalent solids, ionic crystals, and transition  metals. As a result, TMCs often demonstrate extremely high  melting points, considerable hardness. and excellent electric  and thermal conductivity.1 The unique combination of these  desirable properties makes them a kind of promising material  for a wide range of applications including hydrogen evolution   electrocatalysts,2,3 high-temperature solar absorber,4 energy  storage devices, and extreme environment applications,5-7.  Particularly, TMCs, such HfC, ZrC, TiC, and TaC, having  the requisite refractoriness to withstand ultra-high temperatures and pronounced resistance to oxidation simultaneously  are well suited for oxidizing environments at temperatures  exceeding 2000°C. Generally, the oxidation resistance of TMCs is largely  attributed to a carbonaceous oxide interlayer (eg, Hf-O-C,  Zr-O-C, and Ta-O-C), located at the interface between the  external oxide layer and internal carbide (eg, HfC, ZrC, and   J Am Ceram Soc. 2020;00:1-13.   wileyonlinelibrary.com/journal/jace   |   1  © 2020 The American Ceramic Society       \\x0c', '2   |   TaC), acting as the primary oxygen barrier.8-10 Despite carbonaceous oxides have been widely detected in binary systems (eg, HfC, ZrC, TiC, and TaC) since 1967 and profound  research has suggested that the oxygen diffusion coefficient  of carbonaceous oxide (eg, Hf-O-C) is about 30 times lower  than that of corresponding oxide (eg, HfO2),9-12 the essential oxygen barrier mechanism of the carbonaceous oxide has  thus far remained unclear. Recently, newly developed high-entropy ceramics (HECs)  and novel single-phase multicomponent ceramics present a  fountainhead of opportunities for exploring and designing  of better performing TMC compounds.13,14 For  instance,   (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)C,15  (Ti0.2Hf0.2V0.2Nb0.2Ta0.2)C,16  (Zr0.25Nb0.25Ti0.25V0.25)C,17 (Hf0.79Ti0.12Ta0.04Cr0.05)C,18 and   Zr0.8-Ti0.2-C0.74-B0.26 were  fabricated and  revealed  lower  thermal conductivity, higher hardness, and thermal stability  than their original monocarbides.19 These new paradigms  based on early transition metal carbides are potentially useful both structurally and functionally. Especially, researchers  have noted their higher temperature capability and further  enhancement in oxidation and ablation resistance.14,15,18-20  However, the multielement carbonaceous oxide of multicomponent TMCs is rarely studied so far. The aim of present work was to elucidate the oxygen  barrier mechanism of the carbonaceous oxide in TMC compounds. Herein, multiscale investigations into the microarchitecture of the carbonaceous oxide and thermodynamic  analysis were conducted in a multicomponent carbide system (ie, Hf0.5Zr0.3Ti0.2C) under ultra-high temperature oxidation. In order to improve the thermal-shock resistance and  decrease the risk of cracking of the carbide during ablation  test, carbide was incorporated into the C/C composite. The  oxygen barrier mechanism of this carbonaceous oxide interlayer was recently revealed from the aspect of diffusion  mode. Moreover, synergism of the multicomponent system,  particularly, the outward Ti diffusion and its positive role in  self-healing ability, was discussed.  2   2.1   | M E T H O D S | Material preparation  A porous C/C composite with a density of 1.10\\xa0 g\\xa0 cm−3 was  selected as the reinforcement of the multicomponent carbides. Hf-Zr-Ti-C was formed in situ within the C/C matrix via reactive\\xa0 melting\\xa0 infiltration (abbreviated as RMI)  in argon at 2000°C for 2\\xa0 hours. In this step, the composite  sample was placed onto the prepared powders and mixed as  50\\xa0at.% Hf-30\\xa0at.% Zr-20\\xa0at.% Ti. During the RMI process,  the solution of molten hafnium-zirconium-titanium was infiltrated into the porous C/C composites. The molten metal  was drawn along the carbon fibers by capillary forces and   reacted with the previously deposited carbon, forming the  carbon/carbide composites.  2.2   |   Ablation test  The ablation behavior of Hf0.5Zr0.3Ti0.2C was tested by an  oxyacetylene torch device. The inner diameter of the oxyacetylene gun tip was 2\\xa0mm, and the distance between the gun  tip and the specimen was 10\\xa0mm. The specimen, processed  into a size of φ30\\xa0 mm·10\\xa0 mm, was fixed in a water-cooled  copper concave and exposed to flame for 60\\xa0 seconds. The  flow rate and pressure of acetylene were set at 0.696\\xa0 L\\xa0 s−1  and 0.095\\xa0 MPa, respectively, and those of oxygen were,  respectively, 1.960\\xa0 L\\xa0 s−1 and 0.400\\xa0 MPa. During the test,  the highest temperature of the oxyacetylene flame reached  about 2500°C at a distance of 20\\xa0 mm which was measured  by an optical pyrometer. The flame width was about 4-5\\xa0mm,  which was less than the specimen diameter (30\\xa0 mm). The  linear ablation rate (LAR) and mass ablation rate (MAR) can  be calculated by the following formulas:  LAR =  d  i − d  f  t  MAR =  m  i − m f  A ⋅ t  (1)  (2)  where di and df are the initial and final thicknesses measured  at the center of the sample, respectively; mi and mf are the initial and final masses of the sample, respectively; t is the test  duration.  2.3   |   Characterization  The phase composition of the composites was investigated  using a rotating-target X-ray diffraction (XRD) analyzer  (D/max 2550vb\\xa0 +\\xa0 18\\xa0 kW, Rigaku Co.) at a scanning rate  of 2° min-1. The morphology was studied under a scanning  electron microscope (SEM, FEI CO., NOVA Nano230)  combined with an energy dispersive spectroscopy (EDS).  Electron probe microanalysis (EPMA, JEOL CO., Jxa8230)  was used to detect the content of carbon and the distributions  of the major elements in the carbides. High-resolution transmission electron microscope (HRTEM) images and selected  area electron diffraction (SAED) patterns were obtained  with an FEI Tecnai G2 F20 microscope equipped with an  X-FEG electron source. TEM samples were prepared by a  focused ion beam (FIB, FEI Helios Nanolab 600i) using the  in situ lift-out technique on cross-sections of the samples.  Bulk density was measured according to the Archimedes  method. The open porosity was measured using the boiling   YE Et al.    \\x0c', '|   3  water according to the ASTM Standard C20-00. The theoretical density of the carbide was calculated from the mass  of the atoms in a unit cell and the lattice parameters measured from the XRD data. In order to get accurate data, pure  multicomponent carbide powder with identical stoichiometric ratio was fabricated by free pressureless spark plasma  sintering at 1600°C for 15\\xa0 minutes.21 And then XRD patterns were obtained at a scanning rate of 0.5°/min from 5° to  90°. The Rietveld refinement on the XRD patterns was carried out using the general structure analysis system (GSAS)  software.  3   |   R E S U LT S  3.1  | Microstructure and constituents of the  multicomponent carbide  A relative dense ceramic phase was clearly visible in the  cross-sectional SEM image as depicted in Figure\\xa0 1A. The  as-fabricated carbide has a density of 8.65\\xa0 g\\xa0 cm−3, and a  relative density of up to 93% of the theoretical density (ie,  9.28\\xa0g\\xa0cm−3).21 The inhomogeneous distribution of the metal  element in the grains can be observed by EPMA; meanwhile, a compositionally gradient distribution in carbide was   constructed by such elemental inhomogeneity. As shown  in Figure\\xa0 1B, the closer the\\xa0 ceramic-carbon interface, the  higher the content of Hf, while the content of Ti adjacent to  the interface was lower than that in the central part of the ceramic phase. Zr elemental distribution was situated between  Hf and Ti but also assumed a sight tendency of central distribution. Such distribution was further proved by the quantitative analysis conducted by EPMA as shown in Figure\\xa0S1  and Table\\xa0 S1. This can be mainly attributed to the different\\xa0reaction\\xa0trend which is dictated by thermodynamic factor;  HfC has the lowest Gibbs free energy, which means it is thermodynamically favorable for the reaction of Hf and carbon,  while Zr is in the next place followed by Ti. This gradient  distribution might alleviate the mismatch in the coefficient  of thermal expansion (CTE) considering that the CTE of HfC  (7.3\\xa0 ×\\xa0 10−6/°C) is lower than that of ZrC (7.6\\xa0 ×\\xa0 10−6/°C),  TiC (9.5\\xa0×\\xa010−6/°C) and is relatively close to that of the C/C  matrix (3.2-5.7\\xa0×\\xa010−6/°C).22,23 Such a favorable distribution  together with the weakened pyrocarbon interfaces results in  an improvement of thermal shock resistance and decreases  the risk of cracking.24 As expected, strong peaks corresponding to carbide and  carbon-based reinforcement can be observed in the XRD patterns (Figure\\xa01D). The major peaks of the carbide can be indexed to f.c.c. HfC (JCPDS 65-8751), revealing the formation   F I G U R E 1   Microstructure and constituents of the multicomponent carbide Hf0.5Zr0.3Ti0.2C. A, Cross-sectional SEM image of  Hf0.5Zr0.3Ti0.2C, B, the corresponding EPMA mapping of Hf, Zr, Ti, and C elements, C, TEM analysis of the multicomponent carbide and the  corresponding selected area electron diffraction, D, full-scale XRD patterns (blue curve) of Hf0.5Zr0.3Ti0.2C with carbon matrix included. E,  Magnified patterns at (111) and the peak positions of pristine HfC (yellow line), ZrC (green line), and TiC (orange line)  YE Et al.      \\x0c', '4   |   F I G U R E 2   A, SEM image of center region after ablation at 2500°C for 60\\xa0s and B, corresponding EDS elemental map, C, XRD patterns of  Hf0.5Zr0.3Ti0.2C carbide after ablation test at 2500°C, D, magnified patterns with peak positions at (−111) and (111) and peak positions of pristine  HfO2 (orange curve) are given by JCPDS cards 78-0049  of a multicomponent pure f.c.c. solid solution. In comparison  to pristine HfC (JCPDS 65-8751), the multicomponent carbide phase exhibited a small shift of XRD peaks to lower  angle, and its peak positions were between pristine ZrC and  TiC as shown in Figure\\xa01E. This can be ascribed to the substitution of Zr and Ti into the HfC lattice. Given the fact that the  radius of a Zr atom (0.160\\xa0nm) is close to and slightly larger  than that of a Hf atom (0.158\\xa0nm) and a Ti atom (0.146\\xa0nm)  while the difference of lattice constants between HfC, ZrC  and TiC is less than 10%, such substitutional solid solution is  expected to form. Indexing the diffraction patterns of the carbide, as demonstrated in Figure\\xa01C, confirmed that the carbide was in pure  f.c.c. phase; this\\xa0appears\\xa0to\\xa0be\\xa0consistent\\xa0with the XRD result.  The interplanar distance of (200) was measured to be 2.345\\xa0Å  whereas the calculated average value of the (200) interplanar distances of TiC, ZrC, HfC, as received from the JCPDS  files, was 2.314\\xa0Å. Such a difference between the measured  and calculated lattice parameter values further confirms that  the metallic atoms were randomly placed within the metallic  sublattice.25  3.2  | Morphology, architecture, and  phase\\xa0composition\\xa0after oxidation  In this work, the extreme oxidizing scenario up to 2500°C  was established by oxyacetylene ablation test which features  severe oxidation, extreme heat flux and relative short duration. In general, lower LAR and MAR, indicating that oxide  layer expands and increases the weight of exceeding material  loss from ablation, imply\\xa0 minor degradation and better dimensional stability of the tested material. The MAR and LAR  of the multicomponent carbide at 2500°C are −2.5\\xa0 μm\\xa0 s−1  and 1.8\\xa0 mg\\xa0 s−1, respectively. It is noteworthy that the rate  of mass loss of Hf0.5Zr0.3Ti0.2C is 5 times lower than that of  conventional ZrC-based material.26 Such a favorable ablation  performance reflects the effectiveness of protection against  oxidizing environments of the multicomponent carbide and  its oxidation products (ie, multicomponent oxide Hf-Zr- Ti-O) at the macro level. SEM images (Figure\\xa0 2A) reveal that, at 2500°C, oxide  partially melted and formed a relatively dense layer in the  central ablated surface. Figure\\xa0 2C,D displays  the XRD   YE Et al.    \\x0c', '|   5  patterns of the ablated surface. A set of peaks corresponding to the monoclinic HfO2 (JCPDS cards 78-0049) was observed, which suggests that, after ablation, multicomponent  oxide solid solutions, Hf-Zr-Ti-O, were formed through the  oxidation of the multicomponent carbide Hf0.5Zr0.3Ti0.2C.  Compositional uniformity of the multicomponent oxide was  further confirmed by the EDS elemental mapping of ablated  surface (Figure\\xa02B). Figure\\xa03A shows the cross-sectional view at the center of  the sample ablated at 2500°C; the external layer of the oxide  scale which was exposed to oxyacetylene f lame experiencing  severe scouring shows a relatively porous\\xa0structure, while the  inner part remains nearly intact, almost free from the voids  originating from evaporants, suggesting less\\xa0 volatility and  considerable thermal stability of the as-formed multicomponent oxide layer at the micro level. The discrepancy between  the external layer and the inner part resulted from both oxygen  partial pressure (pO2) gradient and the thermal gradient perpendicular to the oxide layer.19,27 As depicted in Figure\\xa03B, A  dark gray\\xa0interlayer can\\xa0be\\xa0observed between the\\xa0oxide layer  and the residual carbide. The EPMA analysis on this combination area confirmed that the interlayer contains oxygen  and carbon which is similar to previously discovered carbonaceous oxide interlayer in the binary carbide system.8-10,28   These carbonaceous oxides, ie, Me-C-O (where Me\\xa0=\\xa0Ti, Zr,  Hf), possess remarkably low diffusion coefficients of oxygen (about 30 times lower than that of outer oxide scale) at  high temperatures.9 Therefore, the carbonaceous oxide plays  a principal role in decelerating oxidative erosion.  3.3  | Microarchitecture of the carbonaceous  oxide region  In order to probe into the essential oxygen barrier mechanism of the discovered carbonaceous oxide interlayer, a thin  slice (Figure\\xa04A) was lifted out from interfacial area cross the  oxide and carbide using FIB and then analyzed using TEM.  Four successive layers with different appearances were documented in Figure\\xa04B,C. As illustrated in Figure\\xa04C, the oxide  layer was composed of the relative coarse oxide grains while  the interlayer consisted of near-equiaxed grains with elongated shape; a relatively uniform sublayer was also observed  between the carbide and the interlayer (Figure\\xa0 4B). In the  following chapters, each layer will be analyzed successively. The HRTEM micrographs provide a clear view of the  oxide layer, as Figure\\xa0 5A,B shows, f ine grain skeletons  were sealed by amorphous phases and these crystalline   F I G U R E 3   A, Cross-section morphology of central ablated region at 2500°C, B, the magnified image of carbide-oxide interface, C,  associated elemental C distribution, D, associated elemental O distribution, E, atom ratios obtained from each of the 14 areas (as shown in  Figure\\xa06B) as a function of distance from the carbide/oxide interface  YE Et al.      \\x0c', '6   |   F I G U R E 4   TEM observation of the thin slice lifted out from the carbide-oxide interfacial region. A, Exact position for lifting out of the  TEM slice B, and C, general view of the carbonaceous oxide interlayer region  F I G U R E 5   TEM observation of oxide layer. A, Representative microstructure of the oxide layer, B, HRTEM image showing amorphous  phase and crystalline oxide, C, diffraction patterns of the crystalline oxide, D, HAADF-STEM image and corresponding EDX results from area B  shown in Figure\\xa0S2. Spots 1 and 2 show the amorphous phase and crystalline oxide, respectively  grains were bonded directly to the amorphous phases. The  indexing of the diffraction patterns from the Four ier transformation of grain skeletons (Figure\\xa0 5C) suggested that  these oxide grains were tetragonal or monoclinic (matches  were obtained with JCPDS 08-0342 and JCPDS 78-0049  respectively). As shown in Figure\\xa0S2, the areas of A, B and C were selected for EDX analysis. Area A was near to the interlayer,   while area C was located at the oxide layer and area B was  distributed between areas A and C. EDX analyses are shown  in Figure\\xa0 5D and Figures\\xa0 S3 and S4 reveals similar results  that the grain skeletons were composed of multicomponent  oxides with lower Ti content, while the carbonand titanium-rich oxides (ie, Ti content higher than 30\\xa0at.%) were the  major constituents of the amorphous phase. More importantly, the carbon content of amorphous phase was found to   YE Et al.    \\x0c', \"|   7  F I G U R E 6   TEM observation of carbonaceous oxide interlayer. A, Representative microstructure of the interlayer, B, TEM image showing  elongated oxide particles with carbon-like tissues surrounding it, C, example of carbon-like tissues indicated by arrow, D, HAADF-STEM image  and corresponding EDX analysis of the elongated oxide particle, E, HAADF-STEM image and corresponding elemental mapping (C, O, Hf, Zr and  Ti)  decrease significantly with an increase in the distance from  the interlayer. The EDX results illustrated in Figure\\xa0S4 suggest  that carbon was consumed significantly and the titanium-rich  multicomponent oxides became a major constituent of the  amorphous phase in the external\\xa0oxide layer. In addition, the  above EPMA analysis (illustrated in Figure\\xa0 3B,E) suggests  that the residue carbon was almost consumed in the slightly  outward region while this outer region remained compact and  consistent with the inner part. Hence, despite the observed  amorphous phases contain a considerable amount of carbon,  titanium-rich multicomponent oxide is still the essential component. It seems that during ablation, small amounts of liquid  phases formed and sealed the defects of the oxide grain skeletons, realizing a self-healing of the oxide layer. Furthermore,  such a self-healing effect at the submicron or nano scales  could substantially promote the densification of the oxide  scale, thereby making it relatively impervious and stable. As Figure\\xa06B,C reveals, the structure of the interlayer can  be distinguished from the oxide layer in terms of its near-equiaxed grains with elongated shape of about 50 -300\\xa0 nm in  length. And they are found to increase in size with the distance from the carbide side to the oxide side (Figure\\xa0 4B,C),  suggesting a coarsening process of oxide particles within the  interlayer. What's more, along the grain boundaries, bright  stripes with typical morphology of carbon were observed in  HRTEM images (Figure\\xa0 6A-C). The elongated grains were  determined to be Hf-Zr-Ti-O oxides from the HAADFSTEM and EDX analyses (Figure\\xa0 6D). Elemental mapping  was carried out to further confirm the composition of these   narrow bright stripes. The results demonstrated in Figure\\xa06E  confirm a carbon-enriched gap area between grains where  bright stripe occupied while the elongated particle mainly  comprising O, Hf Zr, Ti, were\\xa0 carbon deficient; it was thus  suggested these bright stripes were carbon, and elongated  multicomponent oxide particles were surrounded by it. This  in turn demonstrated the microstructure and the constituents  of the interlayer; actually, this carbonaceous oxide interlayer  is made up of elongated multicomponent oxides with carbon  layers surrounding it. Besides, Ti clustering regions occurred  around these oxide grain boundaries were also observed in  elemental mapping (Figure\\xa06E). The HRTEM images in the region adjacent to carbide are  shown in Figure\\xa0 7A,B; a relative uniform sublayer with evident  interface was sandwiched between  the  interlayer and  carbide. Within this sublayer, a discrete distribution of carbon-like tissues among nanocrystalline grains which ranged  from 5 to 20\\xa0 nm in diameter was observed in the HRTEM  image (Figure\\xa0 7C). Here, the EDX mapping was carried out  in the sublayer (Figure\\xa0 S5B,C), and the results showed that  regions of nanocrystalline oxides were\\xa0 carbon deficient while  the rest\\xa0 of\\xa0 the regions were carbon enriched, indicating that  nanocrystalline oxides were surrounded by carbon, further confirming that the discrete tissues shown in Figure\\xa07C were carbon. The SAED patterns (Figure\\xa07A) revealed nanocrystalline  grains to be monoclinic oxides arranged in random orientations.  However the crystal texture of residue carbide remained unaltered; the identification was confirmed by the presence of (2 0 0)   and (2 2 0) carbide spots in the SAED patterns (Figure\\xa0 7A).   YE Et al.      \\x0c\", '8   |   F I G U R E 7   TEM observation of the sublayer. A, Microstructure and electron diffraction patterns of carbide and sublayer, B, TEM image  showing the sublayer and the interlayer, C, HRTEM image of the sublayer  F I G U R E 8   EDX line scan analysis reveals element distributions from carbide to oxide. (A) The positions of the scanning line, (B) lines of  black, blue, red, and green represent Hf, Zr, Ti, and O, respectively, the positions of the scanning line, and the schematic illustrations of the selfhealing mechanism from the view of the interlayer (C) and the oxide layer (D), respectively  Interestingly, the sublayer resembles a scaled-down interlayer  compositionally and structurally; such a resemblance implies  the sublayer may be an infantile period interlayer. Therefore,  the microarchitecture of  the carbonaceous  oxide interlayer region can be described as follows: in the  outer region, oxide grain skeletons sealed by amorphous oxides formed an integral and low-porosity oxide layer. Beneath  it, there was a carbonaceous oxide interlayer composed of  elongated oxide particles about 50 -300\\xa0nm in length with carbon surrounding it, followed by a relatively uniform sublayer   constituted by nanocrystalline oxides with a small amount of  discrete carbon. Moreover, based on the above observation  a reasonable description of the formation and evolution of  the carbonaceous oxide interlayer was also proposed: In the  deepest place of the oxide layer, ie, the reaction interface, due  to the extremely low oxygen activity which was caused by oxygen partial pressure (pO2) gradient,27 oxygen atoms would  first dissolve into the carbide lattice and form f.c.c. oxycarbide.29 Once the pO2 rose in the outer region and ultimately  reached the critical oxygen concentration favorable for phase   YE Et al.    \\x0c', \"transition, the multicomponent carbide was oxidized into  nanocrystalline oxide. Meanwhile, carbon atoms released  from the sublattice formed some nanosized carbon tissues.  And then the nanocrystalline oxides of the sublayer further  developed into near-equiaxed oxides with an elongated shape  possibly due to the pO2 rise. With the growth of sublayer's  nanocrystalline oxide, carbon migrated to the grain boundaries of equiaxed oxides formed lamellar carbon tissues (this  process is documented in Figure\\xa07C and Figure\\xa06C).  4   |   D I S C U S S I O N  |   4.1  Self-healing mechanism of the  multicomponent oxide layer  Generally, a self-healing multi-oxide scale composed of  crystalline oxide skeletons and a liquid phase is desirable for  better protective performance.30 Since the solid grain skeletons provide a framework for the melting phase to be retained  and not removed by shear force, the melting phase could in  turn seal the pores and defects making the inner oxide layer  relatively impervious and stable. In this work, the formation of such advantaged structure  was related directly to the physicochemical synergism of the  designed multi-element system. First, given the incremental concentration of Ti from carbide matrix to oxide scale  whereas the concentration of Hf and Zr significant decreased  (Figure\\xa0 8A,B, to avoid the inf luence of elemental gradient  perpendicular to the interface which was demonstrated in  Figure\\xa0 1, the EDX line scan analysis was set parallel to the  interface), and the clustering tendency of Ti at grain boundaries in the interlayer (see Figure\\xa06D), outward Ti diffusion  was reasonable deduction. It seems that during ablation, some  Ti atoms were released from the solid solution oxides and  diffused outward by short-circuit diffusion in grain boundaries,31 while a similar outward Ti diffusion phenomenon was  also observed by several groups.32,33 Owing to the outward  Ti diffusion, some titanium-rich oxides could be formed at  the grain boundaries while such diffusion also lowered the  Ti content in solid solution oxide skeletons. In the present  work, the content of titanium in our multicomponent system  was designed to be 20\\xa0at.%, from the HfO2-TiO2 and ZrO2- TiO2 phase diagrams as shown in Figure\\xa0 S6. It is observed  that steeply rising liquidus temperature appears near the region containing 20-30\\xa0mol% TiO2. Based on these phase diagrams, it is inferred that the melting phase could formed in  situ in the titanium-rich region due to its lower melting point  while the titanium-deficient particle with a higher melting  point could remain solid and build up an adherent skeleton.  This is further supported by the microanalysis shown in  Figure\\xa05D which indicates titanium contents of glassy oxide  and the crystalline oxide skeleton were determined to be   |   9  33.20% and 10.05%, respectively; Figure\\xa0 6D shows results  consistent with this. Interestingly, the oxide grain boundaries provided the  path for both outward Ti diffusion and elemental carbon releasing (as indicated in Figure\\xa06E) simultaneously; this process is illustrated in Figure\\xa0 8C. Therefore, it is also inferred  that, during ablation, once carbon on the grain boundaries  depleted, some glassy titanium-rich oxide could subsequently  fill these pores and defects in situ (as illustrated in Figure\\xa08D),  making the oxide scale compact and relatively impervious. In  other words, such outward Ti diffusion conferred a self-healing ability at the submicronor nano scales on the oxide layer,  instead of a porous layer. Besides, the low vapor pressures of  the oxides associated with Zr and Hf aided to sustain a stable  and adherent oxide skeleton at high temperatures.5 Additionally, as shown in Figure\\xa0 1B, the gradient distribution of metal elements at the micron scale exhibits a  small quantity of Ti-rich regions formed in local areas at the  center of the ceramic, such as Pos4 shown in Figure\\xa0S1 and  Table\\xa0 S1. Therefore, the Ti-rich oxides on the ablating surface would tend to melt due to their relatively low melting  point according to the phase diagrams (Figure\\xa0 S6), when  the Ti content is higher than 30\\xa0 at.%, while the robust solid  oxides formed in most areas have a higher Hf content. The  melts deriving from the intrinsic Ti-rich oxides (eg, from the  center of the ceramic) were retained by the robust skeleton  of solid oxides, which gives the solid-liquid oxides a good  self-healing ability at the micron scale and a low mass loss.  To a certain extent, such a multicomponent oxide layer could  prevent fast oxygen transport through it and block heat f low,  thereby evidently enhancing integral protection performance.  Furthermore, during ablation this compact oxide layer would  necessarily result in a considerable temperature and an oxygen activity gradient in the direction of scale thickness.19,27  4.2  | Oxygen barrier mechanism of the  carbonaceous oxide  It is known that, during oxidation, pores and defects in the  oxide layer provided the fastest and the straightest diffusion  path for gaseous oxygen diffused to the reaction interface  which would significantly increase the oxidation rate.34,35 In  addition, atomic oxygen diffusion preferentially along grain  boundaries also aggravated the oxidation process,36 while  oxygen diffusion through the inner oxide grain is noticeably  slower, which is about four orders of magnitude slower than  the grain boundary diffusion.37 What's more, the oxidation  rate is generally believed to be controlled by the transport of  oxygen across the oxide film.38 Returning to the diffusion barrier nature of the carbonaceous oxide interlayer, the as-formed carbon may particularly  be associated with the oxygen barrier effect. To deepen the   YE Et al.      \\x0c\", '10   |   understanding of the as-formed carbon, volatility diagrams  of HfC, ZrC and TiC were constructed here according to  the reported method.27,39-41 As a valuable tool for oxidation  behavior study, volatility diagram shows the thermodynamically predicted stable solid phase, the concomitant gaseous  species, and their vapor pressures at different oxygen partial  pressures and temperatures.40 However, these diagrams are  isothermal plots,40 as in the case of the underlying interface, aforementioned temperature gradient across the oxide  layer need to be taken into account. A mathematical model  presented by Guo et al indicates a considerable temperature  gradient (ie, a 500-μm distance is about 900°C) in the ZrO2  scale.42 Given the lower thermal conductivity of HfO2 than  that of ZrO2, the temperature variation among the HfO2dominant scale (ie, Hf-Zr-Ti-O) with a thickness of about  500\\xa0μm (see Figure\\xa03A and Figure\\xa08B) would not be less than  900°C.43 Therefore, 900°C can be deemed as the critical value  of the temperature gradient in the present work; as the test  temperature reached about 2500°C, the thermodynamic temperature of the volatility diagram could be fixed at 1626.85°C  (1900\\xa0K) for the convenience of calculation. Meanwhile due  to a possibly larger temperature gradient resulting from lower  thermal conductivity of HfO2 compared to ZrO2, the volatility diagrams at 1326.85°C (1600\\xa0 K)/1426.85°C (1700\\xa0 K)  /1526.85°C (1800\\xa0K) were calculated. The calculated volatility diagrams of HfC, ZrC, and TiC  at 1626.85°C (1900\\xa0 K) shown in Figure\\xa0 9A-C imply that  as-formed carbon can exist stably at the MC-MO2 (M\\xa0=\\xa0Hf,   Zr) interface where the local oxygen pressure (pO2) is too  low to oxidize this carbon component. That is to say, during  oxidation, carbon is thermodynamically stable within a specific pO2 range and such properties were inherited by the  Hf0.5Zr0.3Ti0.2C solid solution. And, meanwhile, the calculation at 1326.85°C (1600\\xa0K)/1426.85°C (1700\\xa0K)/1526.85°C  (1800\\xa0 K) show that the lower the temperature, the larger  the  thermodynamically stable pO2 range for carbon (see  Figure\\xa09E), indicating that the carbon is more likely to form  when the temperature gradient is larger. These analyses provide additional evidence that the carbon of the carbonaceous  oxide interlayer is expected to form and can exist stably in  TMCs from the aspect of thermodynamics. Hence, as indicated earlier, the compact arrangement of  the near-equiaxed oxide together with a fine nanostructure of  carbon layers filling grain boundaries guaranteed the integrity  and low porosity of the carbonaceous oxide interlayer. Such a  unique structure minimized the fastest gaseous oxygen diffusion. More importantly, the thermodynamic analysis and the  TEM results (Figure\\xa06E) have together revealed that carbon is  able to stably exist along the oxide grain boundaries within the  carbonaceous oxide interlayer. In this context, this carbon could  strongly retard oxygen diffusion along the grain boundaries,  resulting in a much slower oxygen diffusion through the inner  grain of Hf-Zr-Ti-O dominating the oxygen diffusion mode  in carbonaceous oxide interlayer as illustrated in Figure\\xa0 10.  This interpretation is also substantiated by the HAADF-STEM  analysis and the corresponding elemental mapping (Figure\\xa06E)   F I G U R E 9   Calculated volatility diagrams for (A) HfC; (B) ZrC; (C) TiC; and (dD) combined HfC-ZrC-TiC at 1900\\xa0K; (E)  thermodynamically stable pO2 range for carbon at\\xa0different\\xa0temperatures  YE Et al.    \\x0c', \"F I G U R E 1 0   Schematic diagram  depicting oxygen barrier mechanism of  carbonaceous oxide  |   11  which revealed the carbon-enriched grain boundaries were the  areas deficient in oxygen. This result infers that oxygen barely  diffused through these grain boundaries. In addition, the calculated volatility diagrams allow us  to fur ther interpret the outward Ti diffusion dur ing oxidation. According to Wagner's theory of oxidation, the inward  diffusion of O and the outward diffusion of cations across  the oxide film dominate the oxidation process.44 As the  volatility diagram (Figure\\xa09D) suggests, the TiC/TiO2 equilibr ium oxygen pressure is significantly higher than that of  HfC/HfO2 and ZrC/ZrO2, which implies that the outer oxide  layer with a higher local oxygen pressure is thermodynamically prefer red for Ti species to oxidize. This thermodynamic factor may contr ibute a lot to the evident outward  migration tendency of Ti when compared with Hf and Zr  5   |   C O N C L U S I O N S  In this work, the oxygen barrier mechanism of the carbonaceous oxide was studied in a single-phase multicomponent  carbide (ie, Hf0.5Zr0.3Ti0.2C) under oxidizing environment up  to 2500°C which was established by oxyacetylene ablation. It  is noteworthy that the MAR of Hf0.5Zr0.3Ti0.2C is 1.8\\xa0mg\\xa0s−1  is 5 times lower than that of conventional ZrC-based materials, indicating the multicomponent carbide owns a favorable  thermal protection property. Importantly, a carbonaceous  oxide interlayer formed between the external oxide layer and  internal carbide and acted as the primary diffusion barrier  against inward oxygen diffusion. The oxygen barrier mechanism is revealed. Because of  the existence of the compact external oxide layer, beneath  this, the interlayer possesses much lower oxygen activity  and temperature that allow carbon to stably exist. Hence, the   carbonaceous oxide interlayer forms a fine nanostructure of  near-equiaxed oxides with carbon filling the grain boundaries  of oxides. Owing to interlayer's unique and dense nanostructure, gaseous oxygen diffusion through pores was minimized.  More importantly, this as-formed carbon strongly retarded  the fast oxygen diffusion along the grain boundaries of oxides, resulting in a much slower oxygen diffusion through the  inner grain of Hf-Zr-Ti-O dominating the oxygen diffusion  mode in carbonaceous oxide. Additionally, during oxidation,  synergism of the designed multicomponent system, particularly, the outward short-circuit diffusion of Ti, conferred  a self-healing ability on the oxide layer, evidently enhancing the integral protection performance against oxidizing  environments. More broadly, the revelation of oxygen barrier mechanism  of the carbonaceous oxide interlayer with multiple elements  may facilitate the design and optimization of high-entropy  TMCs (eg, high-entropy UHTCs).  AC K N OW L E D G M E N T S  This work was supported by the National Natural Science  Foundation of China (51602349 and 51602351), Fundamental  Research Funds for the Central Universities, Innovation-drive  Project of Central South University, and Key research and development plan of Hunan province (grant no. 2018GK2061).  O RC I D  Yi Zeng\\xa0 Xiang Xiong\\xa0 Ping Xiao\\xa0   https://orcid.org/0000-0002-0899-3786   https://orcid.org/0000-0002-8189-5760   https://orcid.org/0000-0002-6063-3681   R E F E R E N C E S   1. Hwu HH, Chen JG. Surface chemistry of transition metal carbides.  Chem Rev. 2005;105(1):185-212.  YE Et al.      \\x0c\", '12   |    3.    2. Chen WF, Muckerman JT, Fujita E. Recent developments in transition metal carbides and nitrides as hydrogen evolution electrocatalysts. Chem Commun. 2013;49(79):8896-909. Jin H, Chen J, Mao S, Wang Y. Transition metal induced the contraction of tungsten carbide lattice as superior hydrogen evolution  reaction catalyst. 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New insight into the formation and oxygen  barrier mechanism of carbonaceous oxide interlayer in  a multicomponent carbide. J Am Ceram Soc.  2020;00:1-13. https://doi.org/10.1111/jace.17143  YE Et al.      \\x0c']"
},{
  "_id": 134,
  "PDF": "Optical Emission Spectroscopy During Plasmatron Testing of ZrB2-SiC Ultrahigh-Temperature Ceramic Composites.pdf",
  "Text": "['Optical Emission Spectroscopy During Plasmatron Testing of ZrB2-SiC Ultrahigh-Temperature Ceramic Composites  Mickaël Playez∗ and Douglas G. Fletcher†  von Kármán Institute for Fluid Dynamics, 1640 Rhode-Saint-Genèse, Belgium  Jochen Marschall‡  SRI International, Menlo Park, California 94025  and  William G. Fahrenholtz,§ Greg E. Hilmas,¶ and Sumin Zhu∗∗  Missouri University of Science and Technology, Rolla, Missouri 65409  DOI: 10.2514/1.39974  Optical  emission spectroscopy  is used to  investigate  the  oxidation of  a hot-pressed ZrB2 -SiC ultrahigh temperature ceramic composite tested in the 1.2 MW Plasmatron facility at  the von Kármán Institute for Fluid  Dynamics. Time-resolved spectra enable the in situ detection and temporal characterization of electronically excited  B, BO, and BO2 species concentrations directly adjacent to the oxidizing sample surface. The evolution of these boron  species correlates well with the transient formation of a complex multilayer oxide scale containing a silica-rich glassy  outer layer that limits oxide growth.  Nomenclature  a  =  activity mass flow rate, kg \\x01 pressure, Pa heat flux, W \\x01 m\\x002 universal gas constant, 8:314 J \\x01 mol\\x001 \\x01 K\\x001 temperature, K  _m  =  s\\x001  P  =  q  =  R  =  T  =  Subscripts  cw  =  cold wall  dyn  =  dynamic  stat  =  static  w  =  wall  I.  Introduction  U LTRAHIGH-TEMPERATURE ceramic (UHTC) composites based on the diborides ZrB2 and HfB2 in combination with silica-forming refractory materials like SiC and MoSi2 are under  active  study  for  high-temperature  aerospace  applications  on  hypersonic vehicles  [1,2]. Both ZrB2 and HfB2 3380\\x0eC, (3245 and (80 \\x06 40 W \\x01 m\\x001 \\x01 K\\x001 ) respectively) [3] and [4,5], making them potentially enabling for shape-stable leading have  extremely  high melting  points  relatively  high  thermal  conductivities  edge and control surface components for which extreme aerothermal  heating rates lead to the failure of conventional aerospace materials.  In high-temperature oxidizing environments, ZrB2 and HfB2 react with oxygen to form transition metal and boron oxides according to  the reaction \\x85Zr; Hf \\x86B2 \\x87 5=2O2 ! \\x85Zr; Hf \\x86O2 \\x87 B2O3  (1)  Both ZrO2 and HfO2 are refractory materials with melting points above 2500\\x0eC. However, crystalline B2O3 melts at 450\\x0eC, and the softening temperature for amorphous B2O3 glass is reported to lie between 560 and 630\\x0eC [6]. The vapor pressure of liquid B2O3 rises rapidly as temperature increases, reaching an atmospheric pressure boiling point at about 2080\\x0eC [7].  Furnace studies of ZrB2 in air generally reveal parabolic oxidation temperatures below \\x181100\\x0eC, rates at indicating that liquid B2O3 permeates the porous transition metal oxide layer and seals the [8-11]. With increasing  surface,  limiting inward oxygen transport  temperature, volatilization of temperatures above \\x181400\\x0eC, the protective B2O3 liquid leads B2O3 by evaporation is so rapid that the porous transition metal oxide layer may be exposed. When this occurs, the oxide scale no longer  to  more rapid oxidation. At  the loss of  limits oxygen transport effectively and the ZrB2 oxidation rates can become linear. The addition of a silica former improves the oxidation resistance of ZrB2 (and HfB2 ) at temperatures above 1100\\x0eC with the formation of amorphous SiO2 , which has a much lower volatility sealing function [11-15]. At than B2O3 1500\\x0eC, B2O3 has a vapor pressure of about 233 Pa, compared with and performs the same 3 \\x02 10\\x004 Pa for silica [16]. The most-studied diboride-based UHTC composites  contain  between 10 and 30 vol % SiC. When oxidized, these materials form a  multilayer oxide scale with an outer silica-rich glassy layer and an  inner SiC-depleted layer, separated at times by a thin layer of ZrO2 entrained in SiO2 [11,17-20]. The relative thicknesses of the different layers depend on test conditions such as time, temperature,  gas composition, etc. The formation of this complex oxide structure  has been rationalized using thermodynamic calculations and species  transport  arguments  [16,21,22]  supported  by  posttest  sample  Received 23 July 2008;  revision received 7 January 2009; accepted for  publication 18 January 2009. Copyright © 2009 by the American Institute of  Aeronautics and Astronautics, Inc. The U.S. Government has a royalty-free  license  to  exercise  all  rights  under  the  copyright  claimed  herein  for  Governmental purposes. All other rights are reserved by the copyright owner.  Copies of this paper may be made for personal or internal use, on condition  that the copier pay the $10.00 per-copy fee to the Copyright Clearance Center,  Inc., 222 Rosewood Drive, Danvers, MA 01923; include the code 0887-8722/  09 $10.00 in correspondence with the CCC.  ∗Senior Research Engineer, Aeronautics  and Aerospace Department,  Chaussee de Waterloo 72; playez@vki.ac.be Member AIAA. †Past Head, Aeronautics and Aerospace Department; currently, Professor,  Mechanical Engineering, University  of Vermont,  201 Votey Hall,  33  Colchester Avenue, Burlington, VT,  05405;  douglas.fletcher@uvm.edu.  Associate Fellow AIAA. ‡Senior Research Scientist, Molecular Physics Laboratory, 333 Ravens wood Avenue; jochen.marschall@sri.com. Senior Member AIAA. §Professor, Department of Materials Science and Engineering, 223 McNutt  Hall; billf@mst.edu. ¶Professor, Department of Materials Science and Engineering, 223 McNutt  Hall; ghilmas@mst.edu.  ∗∗Ph.D. Candidate, Department of Materials Science and Engineering, 223  McNutt Hall; Sumin.Zhu@us.vesuvius.com.  JOURNAL OF THERMOPHYSICS AND HEAT TRANSFER  Vol. 23, No. 2, April-June 2009  279  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 \\x0c', 'characterization  that  typically  uses  regular  or  field-emission  scanning electron microscopy (FESEM), energy dispersive x-ray  spectroscopy (EDS), and x-ray diffraction.  The  oxidation  of  hot-pressed -30SiC) was  ZrB2 recently studied in the high composites  containing  30 vol % SiC (ZrB2 temperature, low-pressure dissociated air flow generated by the  1.2 MW Plasmatron facility of the von Kármán Institute for Fluid  Dynamics  (VKI). Plasmatron and arcjet  facilities generate high enthalpy,  highly  dissociated  flows  that  better  simulate  the  thermochemical environment expected in service on a hypersonic  vehicle  than do furnace  environments. Results describing oxide  surface properties, composition, and microstructure as a function of  test  time  and Plasmatron conditions  are presented in a  separate  publication [23]. Here we report emission spectroscopic measure ments performed as an adjunct  to this test series. These measure ments  were  prompted  by  the  observation from BO2 \\x85A2\\x05u ! X2\\x05g \\x86 of strong green to UHTC specimen surfaces during  fluorescence,  suspected  to  originate  emission,  directly  adjacent  exposure to the Plasmatron gas flow. Spectral emission data confirm  this  hypothesis  and, system and the B\\x852 S1=2 ! 2Po additionally, capture emission Time-resolved emission measurements can  from the  BO\\x85A2\\x05 ! X2\\x06\\x87 \\x86  1=2;3=2 \\x86  transi tions.  follow  the  temporal evolution of emission intensities from these electronically  excited boron species as the UHTC sample oxidizes. Because all  boron species in the gas phase must originate from the volatilization  of B2O3 , such time-resolved spectra provide an in situ monitor for UHTC oxidation and independent data for correlation with oxide  growth models.  II.  Experiment  A.  Test Specimens  UHTC specimens were hot pressed as thin disks, starting with  commercial powders (H.C. Starck grade B ZrB2 and grade UF-10 SiC) that were attrition milled using tungsten carbide media for two  hours and dried by rotary evaporation. The dried powders were hot  pressed in boron-nitride-coated graphite dies lined with graphite foil. The hot press was heated under vacuum (\\x1820 Pa) at \\x1820\\x0eC \\x01 min\\x001 to 1450\\x0eC, held for 1 h, then heated at the same rate to 1650\\x0eC and  argon and heated at \\x1820\\x0eC \\x01 min\\x001 held for another hour. The hot press was then backfilled to 1 atm with to a temperature of 1900\\x0eC, at which a pressure of 32 MPa was applied and the specimens held for  \\x1820\\x0eC \\x01 min\\x001 to room temperature and, at \\x181750\\x0eC, the load was 45 min. At the end of this hold time, the hot press was cooled at removed. The pressed samples were approximately 3-4 mm thick  with diameters of about 32 mm. The large faces of the hot-pressed  disks were diamond ground flat and a 30 deg bevel was diamond  ground into the disk edge so that  the samples would fit  into the  standard ESA 50 mm stagnation  point  test fixture  used  in  the  Plasmatron facility.  Posttest  samples were  prepared  for microscopy  by  cutting  specimens perpendicular  to their oxidized faces and polishing the  cross  sections  to  a  0:25 \\x16m finish  using  diamond  abrasives.  Microstructure  and  composition  characterization was  performed  using high-resolution field-emission scanning electron microscopy  with  a  JEOL 6100  instrument  (JEOL, Ltd., Tokyo,  Japan)  in  conjunction with energy dispersive x-ray spectroscopy (EDAX,  Mahwah, New Jersey).  B.  Plasmatron Conditions  Plasma oxidation experiments were performed in air. The 1.2 MW  Plasmatron facility at  the von Kármán Institute [24,25] generates a  high-enthalpy  subsonic gas flow using inductive  coupling. The  plasma flow is directed through a quartz tube into a vacuum chamber  in which test models and probes for measuring dynamic pressure and  stagnation point heating rates are mounted on water-cooled arms that  can be swung into and out of  the flow. Test conditions are set by  adjusting the generator power; the air mass flow rate, _m; the static pressure in the vacuum chamber, Pstat ; and the exposure time in the flow. The cold-wall stagnation point heat flux, qcw , and the dynamic  pressure, Pdyn , are measured using water-cooled probes the same size and external shape as the test specimen holder so that important flow  parameters  (dynamic pressure, velocity gradient) are reproduced.  Sample  surface  temperatures  are measured with  a  two-color  pyrometer  (Marathon Series MR1SC, Raytek Corporation, Santa  Cruz, California) at an acquisition rate of 1 Hz through a glass  window at an incident angle 29 deg off-normal.  The data are presented for  two UHTC specimens  tested under  nominally identical test conditions; see Table 1. The boundary-layer edge enthalpy, gas 25 kJ \\x01 g\\x001 , temperature, and velocity are estimated to be 130 m \\x01 5800\\x0eC, s\\x001 computational fluid dynamics calculations performed at VKI (details  approximately  and  based  on  of these computations are provided in a separate manuscript [23]).  Under these conditions, molecular oxygen is completely dissociated  and the number density of atomic oxygen and atomic nitrogen are  approximately equal.  Sample  17 was  used  in  preliminary  calibration  experiments  associated with the earlier UHTC test series, during which it was  rates of 8 and 16 g \\x01 exposed to air plasma flows under various test conditions (mass flow s\\x001 , power 200 kW, surface temperatures ranging from 1260 to 1400\\x0eC) levels varying between 150 and  for  many (at  least 10) minutes. As a result,  the specimen surface was  preoxidized, with a multilayer oxide structure as described in the  Introduction. Sample 7 was a virgin specimen with no surface oxide  layer present.  C.  Emission Measurements  The emission from the hot gas in front of the test specimen was  collected using the optical system shown in Fig. 1. The collection  optics were aligned such that the optical axis was perpendicular to the  plasma jet and parallel to the sample face along its symmetry plane,  with the 0.4-mm-diam light collection volume tangent to the sample  surface. The light emitted by the gas  is collected by a spherical  mirror, which images the emission, after a second reflection off a flat  mirror and onto the entrance of an optical fiber. The fiber transmits  the light to an Ocean Optics HR4000CG-UV-NIR spectrometer for  the dispersion and recording of the emission spectrum.  This high-throughput spectrometer (f=4) images emission onto a  3648-element linear Si charge-coupled device (CCD) array over the 200-1100 nm wavelength range with a 300 groove=mm grating and  a fixed 5 \\x16m slit width, which provides a spectral  resolution of  0.25 nm. The optical system was calibrated using a tungsten ribbon  lamp (OSRAM ref. WI 17G). The spectral response of the system was determined for the wavelength interval of 350-800 nm, during  which the calibration signal  is sufficiently strong to be measured.  Measured spectra were collected with a time resolution of 1 s.  III.  Experimental Results  Pictures of the sample holder in the plasma flow were taken every  second. Figure 2 compares  two pictures  taken 55 s after  sample  injection into the plasma stream; preoxidized sample 17 is on the left  (Fig. 2a) and virgin sample 7 is on the right (Fig. 2b). The difference  in the color of  the gas surrounding the sample holder  in the two  images is obvious. A strong green fluorescence is seen clearly in  Fig. 2b, but not in Fig. 2a. A similar green emission is well known in the field of boron combustion [26-29]. A small amount of green light  was also observed at earlier times during the testing of sample 17, but  it disappeared quickly. Sample 17 was tested before sample 7, and  the same SiC cover was used for both experiments, indicating that the  SiC cover does not produce this emission phenomenon  power \\x88 220 kW , Pstat \\x88 104 Pa, and _m \\x88 16 g \\x01 s\\x001 Table 1 Test conditions, where generator for all tests  Sample  qcw , W cm\\x002  Pdyn , Pa  Tw , \\x0eC  Time, min  17 (preoxidized)  120  39  1616  3:20  7 (virgin)  119  38  1602  5:10  280  PLAYEZ ET AL.  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 \\x0c', 'PLAYEZ ET AL.  281  Fig. 1  Plasma emission collection system.  Fig. 2  Samples during the Plasmatron testing 55 s after injection of the probe: a) sample 17, and b) sample 7.  Figures 3a and 3b show emission spectra collected 55 s after the  injection of samples 17 and 7 into the plasma stream, corresponding  in time to the photographs in Fig. 2. Figure 3a shows that emission originating from the doublet B \\x852 S1=2 ! 2Po present in the spectrum collected in front of sample 7 but not sample  1=2;3=2 \\x86  transitions  is  17. The wavelength region below 350 nm is not calibrated, and the  spectra in Fig. 3a are presented as raw signal features associated with emission in the BO2 \\x85A2\\x05u ! in arbitrary units. In X2\\x05g \\x86 Fig. 3b, and the BO \\x85A2\\x05 ! X2\\x06\\x87 \\x86 systems are evident in the spectrum collected in front of sample 7, but are absent in the spectrum  collected in front of sample 17. Both spectra in Fig. 3b show strong g \\x86 first negative system in the emissions from the N\\x87 350-450 nm wavelength region. The N first negative emission  \\x85B2\\x06\\x87 u ! X2\\x06\\x87  \\x87  2  2  originates from the plasma freestream. Both spectra also show sharp  emission lines near 589 and 546.7 nm. The former feature has a much  larger intensity adjacent  to the surface of sample 7 than sample 17  and is not present in the plasma freestream; it can be assigned to the strong Na \\x852Po 1=2;3=2 ! 2 S1=2 \\x86 doublet emission [30]. Further testing has shown that the apparent 546.7 nm emission feature results from a  failed pixel in the linear CCD array.  Preliminary computations of the radiation emitted by the BO \\x85A2\\x05 ! X2\\x06\\x87 \\x86 system were performed. The emission of the two BO isotopes, 11BO and 10BO, was modeled independently using  the spectroscopic constants of Mélen et al. [31]. Constants for states with vibrational quantum numbers \\x1d0 varying from 0 to 8 and \\x1d00 from 0 to 8 for 11BO and \\x1d0 from 0 to 8 and \\x1d00 from 0 to 6 for 10BO are  available in this  reference. These constants were extrapolated to  vibrational quantum numbers up to 10 for the A state and up to 17 for  the X state [32]. Figure 4 shows a comparison of a model BO \\x85A2\\x05 ! X2\\x06\\x87 \\x86 spectrum with the difference of the spectra measured in front of the  two samples  (i.e.,  the  spectrum of  sample 7 minus  spectrum of  sample 17) 55 s after injection into the plasma stream. Owing to the  observation that the recorded spectra from the preoxidized and virgin  samples agreed very well  in regions above 750 nm in which boron  species emission was absent, an unadjusted subtraction was made.  This subtraction effectively removes the nitrogen emission features  from the  spectrum. The model  spectrum was  computed  for  a  temperature of 6500 K using the hypotheses of local thermodynamic  equilibrium emission and an optically thin medium. The two spectra  were normalized individually by their peak values and offset  for  clarity.  Although  the  computed  and  measured  spectral  intensity  distributions do not match well,  the band structure is similar in the  two spectra, particularly in the short wavelength region in which  overlap with BO2 features (whose bandheads are indicated in Fig. 4) is not important. The observed difference between the two intensity  distributions  could be  caused by temperature  and concentration  nonuniformities in the flow, self-absorption along the line of sight,  uncertainties in the spectroscopic constants used in the simulation,  and/or the overlap of experimental BO features with other emission  features, such as those of BO2 . Spectra computed for temperatures lower than 6500 K show worse agreement with the experimental  intensity  distribution  in  the  short wavelength  region, whereas  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 \\x0c', 'spectra computed for higher  temperatures  show somewhat better  agreement.  Based on computational fluid dynamics simulations performed at  \\x186100 K [23]. However, the cylindrical emission collection volume VKI, the plasma freestream temperature is estimated to be about is tangent to the \\x181725 K sample surface (55 s into the test run) and the gas temperature in this part of the boundary layer is much lower,  reaching only about 3000 K at a distance of 0.5 mm off the surface.  The prominent short wavelength features in the experimental spectra nonequilibrium “hot”  suggests  the  possibility  of  a  vibrational  distribution in the electronically excited BO. More data analysis and  modeling are  required before  this possibility can be  confirmed.  Nevertheless,  the comparison made in Fig. 4 underlines the strong  contribution of BO emission to the observed fluorescence.  Figure 5 compares the transient surface temperatures measured on  samples 7 and 17 with the temporal emission signatures of BO2 (518.8 nm), BO (404.1 nm), and B (249.9 nm). These three  wavelengths were chosen to minimize overlapping emission from  multiple species. The pyrometer registers surface temperatures above 1100\\x0eC; the first portion of the temperature trace for sample 17  was  lost  during  data  acquisition. For  sample  7,  boron  species  emission signatures exceeds \\x181250\\x0eC, about 40 s after sample injection into the stream. rise very rapidly as the surface temperature A quasi-steady surface temperature around 1600\\x0eC is attained about  80 s after  sample injection, and all  three boron species emission  signatures decay steadily from this time onward. The calibrated BO2 and BO emission intensities peak at intermediate times around 55 s.  Although sample 17 experiences the same nominal heating condition  resulting in similar surface temperatures, the boron emission features  between 40 and 140 s are absent. After about 160 s,  the emission  levels  in front of both samples become comparable. The test of  sample 17 was ended early as no further evolution of the emission  spectrum for the preoxidized sample was observed.  IV.  Discussion  Thermodynamic  data  from [33] were  used  to  calculate  the  pressures of gaseous species in equilibrium with liquid B2O3 at 1600\\x0eC to determine the species that are likely to volatilize from the  specimen surface. Potential reactions among gaseous species were  monatomic oxygen at a pressure of 3 \\x02 103 Pa (PO \\x88 3 \\x02 103 Pa or ignored for this analysis. The atmosphere was assumed to be 0.02961 atm). Data for the standard Gibbs free energies of formation (\\x01Go f ) were extracted from the tables and used to calculate changes in the standard Gibbs free energy of reaction (\\x01Go rxn ), which were subsequently converted to pressures of the gaseous species. The volatilization reaction B2O3\\x85l\\x86 \\x87 O\\x85g\\x86 ! 2BO2\\x85g\\x86 ered as an example. At 1600\\x0eC, \\x01Go is 64:26 kJ \\x01 mol\\x001 and the is consid(PBO2 ) was determined to be 2515 Pa  rxn  equilibrium pressure of BO2 from the relationship  \\x01Go  rxn \\x88 \\x00RT ln  \\x12  P2  BO2  aB2 O3  PO  \\x13  (2)  (aB2 O3 \\x88 1). Similar calculations assuming unit activity for B2O3 were performed for other boron species and the values are  summarized in Table 2. Based on these calculations, the flux of boron  species  leaving  the  sample  surface  during  the  initial  stages  of  oxidation in the Plasmatron should be predominantly BO2 and B2O3 , with much smaller contributions from BO and B2O2 and only trace amounts of B, B2O, and B2 . Emission spectroscopy detects the presence of electronically excited species. The BO2 \\x85A2\\x05u \\x86 [34], BO\\x85A2\\x05\\x86 [35], and B\\x852 S1=2 \\x86 [30] excited states are located about 2.3, 3.0, and 5.0 eV above their  respective  ground  states.  It  seems  unlikely  that  any  of  these  electronically excited boron species would volatilize from liquid  B2O3 observed BO2 , BO, therefore result from some combination of gas-phase chemistry and  directly. The  and B emissions must  energy transfer between boron volatiles  in the ground state  and  freestream species. Based on the presence of  large atomic oxygen  and nitrogen concentrations at the sample surface, the observation of  0 200  220  240  260  280  300  100  200  300  400  500  600  700  800  B (2S1/2-2P1/2,3/2 )   Sample 7  Sample 17  R  A  W  S  I  G  N  A  L  ,  a  .  u  .  WAVELENGTH, nm  350  400  450  500  550  600  650  0  500  1000  1500  2000  2500  3000  Na  N2 + Emission  BO2(A-X )  BO(A-X)  I  N  T  E  N  S  I  T  Y  ,  W  ⋅  c  m   2   ⋅  s  r   1  ⋅  µ  m   1  WAVELENGTH, nm   Sample 7  Sample 17  a)  b)  Fig. 3 Spectra collected 55 s after sample injection showing strong B, BO, and BO2 emissions for sample 7 and the absence of boron emission signatures for sample 17: a) raw signal, and b) calibrated intensity.  Boron oxide bandhead positions are from Spalding et al. [26].  0.0 350  400  450  500  550  600  650  0.5  1.0  1.5  2.0  BO (A-X)  N  O  R  M  A  L  I  Z  E  D  I  N  T  E  N  S  I  T  Y,  a  .  u  .  WAVELENGTH, nm   Difference Spectrum  Model BO (A-X ) Spectrum  BO2 (A-X)  Fig. 4  Comparison of  the  computed (6500 K) BO \\x85A2\\x05 ! X2\\x06\\x87 \\x86  emission spectrum with the difference of the spectra measured in front of  the two samples (i.e., the spectrum of sample 7 minus the spectrum of  sample 17) 55 s after injection of the probe. The spectra are normalized  independently by their peak values and are offset along the intensity axis  for clarity. The red lines indicate the alignment of computed BO features  with emission features  in the the BO2 \\x85A2\\x05u ! X 2\\x05g \\x86 experimental difference  spectrum;  the  emission bandheads  for  system are  from  Spalding et al. [26].  282  PLAYEZ ET AL.  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974           \\x0c', 'boron atom emission at  relatively low gas  temperatures near  the  surface,  and  BO spectra  that  suggest  nonthermal  vibrational  distributions,  the boundary-layer gases immediately adjacent to the  surface  are  probably  not  in  thermochemical  equilibrium.  The  observed  electronically  excited  boron  species  emissions may  therefore contain significant contributions from kinetically driven  chemiluminescent  reactions. The extent of  thermal and chemical  nonequilibrium in the emitting species cannot be evaluated without a  coupled fluid dynamics and finite-rate chemistry computation of the  plasma-generated  species  with  the  volatile  boron  species,  a  complicated problem beyond the scope of this paper. The reaction of gas-phase B2O3 with atomic oxygen, B2O3\\x87 O ! 2BO2 , is thermodynamically favored (negative \\x01Go rxn ) for all test temperatures [7]. The Gibbs free energy for the thermal  dissociation of the larger boron oxide molecules becomes favorable [7]: B2O ! BO \\x87 B at \\x182915\\x0eC, B2O2 ! 2BO at \\x183030\\x0eC, B2O3 ! BO \\x87 BO2 at \\x183380\\x0eC, and with increasing temperature BO2 ! BO \\x87 O at \\x184015\\x0eC. Dissociation of these species is additionally promoted by collisions with highly energetic electroni cally and vibrationally excited nitrogen molecules produced by the  Plasmatron. Atomic oxygen reactions with the smaller boron species  are the most likely candidates for generating chemiluminscence. At 3000\\x0eC,  the  following  reactions  are  both  thermodynamically  produce the BO2 \\x85A2\\x05u \\x86 and BO\\x85A2\\x05\\x86 states [7]: B \\x87 O \\x87 M ! favored (negative Gibbs free energy) and sufficiently exothermic to BO \\x87 M \\x87 8:53 eV, B2 \\x87 O ! BO \\x87 B \\x87 5:45 eV, B2O \\x87 O ! 2BO \\x87 4:47 eV, BO \\x87 O \\x87 M ! BO2 \\x87 M \\x87 5:79 eV. and (Here, M stands for the third body collision partner.)  The temporal evolutions of the boron species emission signatures  shown in Fig. 5 are consistent with the general understanding of  ZrB2 the reaction sequence  -SiC oxidation observed at different temperatures [11,17] and formation at 1500\\x0eC  for multilayer oxide  described by Fahrenholtz [16]. The UHTC sample temperature rises  very rapidly after 800\\x0eC, both ZrB2 1200\\x0eC temperature and SiC oxidation is negligible.  injection into the plasma stream. Below about In the 800-  interval,  the ZrB2 whereas the SiC oxidation rate remains insignificant,  oxidation  rate  increases  leading to a  surface oxide composed of ZrO2 and SiC grains in a glassy B2O3 matrix. As the sample temperature rises through this regime, the  B2O3 amounts of boron oxide species temperature surpasses 1200\\x0eC,  volatilization  rate  also  accelerates,  injecting  significant  into the boundary layer. As  the  the oxidation rate of SiC becomes  significant and the outer scale transforms into a borosilicate glass.  Further increases in temperature raise the rates of both silicon oxide  production and boron oxide loss, leading to the steady depletion of  boron from the glassy scale. The volatilization rate of silicon oxide is  many orders of magnitude lower  than that of boron oxide in this  temperature  range.  The  resulting  silica-rich  outer  oxide  layer  becomes an effective oxygen diffusion barrier and is responsible for  the  parabolic  oxidation  kinetics observed during steady-state temperatures near 1500\\x0eC [16]. As  oxidation experiments at  the  silica-rich layer  thickens, oxygen diffusion limitations  lower  the  partial pressure of oxygen at the internal reaction interface, slowing  production of B2O3 . At sufficiently low oxygen partial pressures, the oxidation of ZrB2 becomes insignificant relative to the (active) oxidation of SiC [16]. The flux of boron into the boundary layer  decreases and eventually ceases, as boron is depleted from the outer  glassy scale and is no longer replaced by B2O3 from in-depth ZrB2 oxidation.  Figure 6 shows a cross-sectional FESEM micrograph of sample 7  after Plasmatron testing, together with corresponding EDS maps of  elemental boron, zirconium, oxygen, and silicon. The EDS maps  show the detection of each element as a white signal on a black  background. Above the virgin material, the FESEM image reveals a \\x1830-\\x16m-thick sublayer of oxidized material uneven, glassy surface layer. In the virgin material,  covered by a  thin,  the EDS maps  show the presence of boron, zirconium, and silicon and the relative -30SiC  absence  of  oxygen,  as  expected  for  unoxidized  ZrB2 material. The EDS maps show both zirconium and oxygen in the  oxidized sublayer, indicating the oxidation of ZrB2 to ZrO2 and the presence of silicon and oxygen in the glassy outer layer  corresponding to the formation of SiO2 SiC. The weakness of the EDS boron signal  through the oxidation of  in both the oxidized  sublayer and the outer glassy layer clearly indicate diminished boron  0  50  100  150  200  250  300  350  1100  1300  1500  1700  BO2   = 518.8 nm  B  R  A  W  S  I  G  N  A  L  ,  a  .  u  .   Sample 7   Sample 17   = 249.9 nm  TIME, s  T  E  M  E P  R  A  T  R U  E  ,  °  C  BO  I  N  T  E  N  S  I  T  Y,  a  .  u  .   = 404.1 nm  Fig. 5  Transient profiles of the surface temperature and the emission signatures of BO2 , BO, and B at 518.8, 404.1, and 249.9 nm, respectively.  Table 2 Calculated pressures of volatile boron species in equilibrium with liquid B2O3 at 1600\\x0eC  Species  Pressure, Pa  BO2 B2O3 BO  2215  1:1 \\x02 10\\x003 917 3:6 \\x02 10\\x007 2:3 \\x02 10\\x0018 1:9 \\x02 10\\x0020 1:3 \\x02 10\\x0038  B2O2 B  B2O B2  PLAYEZ ET AL.  283  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 λ       λ       λ \\x0c', 'concentrations relative to the virgin material and correspond well  with the scenario of boron oxide volatilization outlined earlier. A surface temperature of \\x181600\\x0eC is still moderate compared with the ambitious goal of sustained operation of UHTC components at 2000\\x0eC and higher [1,2]. As temperatures approach 2000\\x0eC, a  similar volatilization of silicon species occurs and the outer silica rich glassy layer recedes, leaving a porous zirconia (or hafnia in the  case of HfB2 -based material) oxide scale [20,21,36]. interesting to explore if similar spectroscopic emission diagnostics  It would be  could be used to detect and track Si and SiO volatilization under such  conditions.  V.  Conclusions  The  emission measurements  presented  here  demonstrate  that  boron species can be detected in the gas phase at Emissions originating from the BO2 \\x85A2\\x05u ! X2\\x05g \\x86 system, ZrB2 -30SiC UHTC material exposed to an air plasma flow. BO \\x85A2\\x05 ! X2\\x06\\x87 \\x86 system, and the B \\x852 S1=2 ! 2Po 1=2;3=2 \\x86 doublet the transitions were identified in spectra collected over the 200-800 nm  the surface of a  wavelength  range. Because  all  gas-phase  boride  species must  ultimately originate from the sample, time-resolved emission spectra  from electronically excited B, BO, and BO2 molecules provide an in situ monitor for the oxidation of the ZrB2 -30SiC UHTC material. The rise and fall of these boron species emissions during the test  series  correlate well with  theoretical models  of  the  transient  formation of complex oxide layers based on competing oxidation,  diffusion, and vaporization processes. Posttest microscopy of  the  oxidized sample confirms the expected depletion of boron in the  oxide scale sublayer beneath the outer silica-rich glass.  Because most UHTC materials  (monolithic or  coating) under  development  for  leading-edge and control surface applications on  hypersonic vehicles contain some boride constituents [1,2], emission  spectroscopy is potentially broadly useful  for  following UHTC  stability and performance in extreme oxidizing Plasmatron and arcjet  test environments.  Acknowledgments  This research was supported by the Ceramics and Nonmetallic  Materials Program of the U.S. Air Force Office of Scientific Research  through contracts F49550-05-C-0020 (Marschall) and FA9550-06 0125 (Fahrenholtz and Hilmas),  the National Science Foundation  through  grants DMR-0435856  (Marschall)  and DMR-0346800  (Fahrenholtz  and Zhu),  and  the European Office  of Aerospace  Research and Development  through contract FA8655-06-1-3078  (Playez and Fletcher). The authors thank Dušan Pejaković of SRI for  performing the FESEM and EDS measurements and Jan Thömel of  VKI for computational estimates of the Plasmatron test conditions.  References  [1] Fuller, J., Blum, Y., and Marschall, J., “Topical Issue on Ultra-HighTemperature Ceramics,” Journal of Vol. 91, No. 5, 2008, pp. 1397-1502.  the American Ceramic Society,  doi:10.1111/j.1551-2916.2008.02481.x [2] Fuller, J., and Sacks, M., “Special Section: Ultra-High Temperature Ceramics,” pp. 5885-6066.  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[36] Bronson, A., and Chessa, J., “An Evaluation of Vaporization Rates of  SiO2  and TiO2  as Protective Coatings  for Ultrahigh Temperature  Ceramics,” Journal of the American Ceramic Society, Vol. 91, No. 5,  2008, pp. 1448-1452.  doi:10.1111/j.1551-2916.2008.02286.x  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 \\x0c', 'This article has been cited by:  1. Jason D. White, Jochen Marschall, Richard A. Copeland. 2013. First measurement of the BO B 2Σ+ radiative lifetime by laser-induced fluorescence via the B 2Σ+-X 2Σ+ transition. Chemical Physics Letters 577, 16-21. [CrossRef] 2. Bernd Helber, Olivier Chazot, T. Magin, Annick HubinSpace and Time-Resolved Emission Spectroscopy of Carbon Phenolic Ablation in Air and Nitrogen Plasmas . [Citation] [PDF] [PDF Plus] 3. Jochen Marschall, Dušan Pejakovic, William G. Fahrenholtz, Greg E. Hilmas, Francesco Panerai, Olivier Chazot. 2012. Temperature Jump Phenomenon During Plasmatron Testing of ZrB2-SiC Ultrahigh-Temperature Ceramics. Journal of Thermophysics and Heat Transfer 26:4, 559-572. [Citation] [PDF] [PDF Plus] 4. Jochen Marschall, Du�an A. Pejakovi�, William G. Fahrenholtz, Greg E. Hilmas, Francesco Panerai, Olivier Chazot. 2012. Temperature Jump Phenomenon During Plasmatron Testing of ZrB 2 SiC Ultrahigh-Temperature Ceramics. Journal  of Thermophysics and Heat Transfer 26:4, 559-572. [CrossRef]  5. Francesco Panerai, Olivier Chazot. 2012. Characterization of gas/surface interactions for ceramic matrix composites in high enthalpy, low pressure air flow. Materials Chemistry and Physics 134:2-3, 597-607. [CrossRef] 6. D Le Quang, Y Babou, P Andre. 2012. Investigations of a microwave plasma source operating with air, N 2 , CO 2 and  argon gases. IOP Conference Series: Materials Science and Engineering 29, 012009. [CrossRef]  7. Jochen Marschall, Douglas G. Fletcher. 2010. High-enthalpy test environments, flow modeling and in situ diagnostics for characterizing ultra-high temperature ceramics. Journal of the European Ceramic Society 30:11, 2323-2336. [CrossRef] 8. Jochen Marschall, Dusan A. Pejakovic, William G. Fahrenholtz, Greg E. Hilmas, Sumin Zhu, Jerr y Ridge, Douglas G. Fletcher, Cem O. Asma, Jan Thoemel. 2009. Oxidation of ZrB2-SiC Ultrahigh-Temperature Ceramic Composites in Dissociated Air. Journal of Thermophysics and Heat Transfer 23:2, 267-278. [Citation] [PDF] [PDF Plus]  Downloaded by SIMON FRASER UNIVERSITY on July 16, 2013 | http://arc.aiaa.org | DOI: 10.2514/1.39974 \\x0c']"
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  "_id": 135,
  "PDF": "Oxidation and its effect on flexural strength of hot pressed ZrB2 SiC composites with VC additives.pdf",
  "Text": "['Journal of Asian Ceramic Societies  ISSN: (Print) (Online) Journal homepage: https://www.tandfonline.com/loi/tace20  Oxidation and its effect on flexural strength of hotpressed ZrB2-SiC composites with VC additives  Shuqi Guo  To cite this article: Shuqi Guo (2020): Oxidation and its effect on flexural strength of hotpressed ZrB2-SiC composites with VC additives, Journal of Asian Ceramic Societies, DOI: 10.1080/21870764.2020.1840701  To link to this article:  https://doi.org/10.1080/21870764.2020.1840701  © 2020 The Author(s). Published by Informa UK Limited, trading as Taylor & Francis Group on behalf of The Korean Ceramic Society and The Ceramic Society of Japan.  Published online: 05 Nov 2020.  Submit your article to this journal   Article views: 81  View related articles   View Crossmark data  Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=tace20  \\x0c', 'JOURNAL OF ASIAN CERAMIC SOCIETIES             https://doi.org/10.1080/21870764.2020.1840701  FULL LENGTH ARTICLE  Oxidation and its effect on flexural strength of hot-pressed ZrB2-SiC  composites with VC additives  Shuqi Guo  Research Center for Structural Materials, National Institute for Materials Science, Tsukuba, Japan  ABSTRACT  The oxidation and its effect on room-temperature flexural strength were investigated in the  hot-pressed ZrB2-SiC composites with 3−7 wt% VC additives exposed to air within the temperature range of 1100−1300ºC, for up to 100 h. The parabolic weight gain versus oxidation  time was measured for the period of 10−100 h. The apparent oxidation activation energy was  determined to be 70−180 kJ/mol, dependent on the amount of VC. The oxidation products  consisted of m-ZrO2, V2O5, and ZrSiO4 at 1100ºC, m-ZrO2, SiO2 (αor β-cristobalite), ZrV2O7, and  ZrSiO4 at 1200ºC and 1300ºC. Moreover, the room-temperature flexural strength of the postoxidized composites varied depending on the amount of VC as well as on oxidation temperature. The change in the strength of the postoxidized composites was associated with the  formation of either outermost dense glassy scale or new defects in the oxide layer during the  oxidation exposure.  ARTICLE HISTORY   Received 30 June 2020   Accepted 19 October 2020   KEYWORDS   ZrB2; SiC; VC additive;  oxidation; flexural strength  1.  Introduction  Zirconium diboride  (ZrB2)  is a refractory transition metal diboride composed of elements from the fourth  to sixth groups of the periodic table. Most of these  diborides have melting points greater than 3000°C,  high thermal and electrical conductivities, chemical  inertness against molten metals, and good thermal  shock resistance, making them potential candidates  for several high-temperature structural applications  [1-3]. Two major problems of ZrB2 ceramics are the  following: (i) poor sinterability because of strong covalent bonds and  low self-diffusivity  [4], and  (ii)  low  oxidation resistance in the air or oxidized atmosphere  at elevated temperature [5,6]. The composite approach  has been successfully adopted to improve the densification and the oxidation resistance of monolithic ZrB2  ceramics. It is known that the addition of SiC to ZrB2  results  in a composite with  improved sinterability,  great strength, and better oxidation resistance [7-12],  and  recently  this composite has become a strong  potential candidate for a variety of high-temperature  structural applications [1-3]. In addition, previous studies of ZrB2-based ceramics showed that groups IV-VI  transition metal carbides could be used as a sintering  activator to aid the densification of ZrB2-SiC composites at a lower temperature [13,14]. Zou et al. [15] had  reported that highly dense ZrB2-SiC ceramics with VC  additions could be fabricated by pressureless at or  above 2000°C for 2 h in Ar. Furthermore, they showed  that density exceeding 99% was obtained at 1900°C for  1 h under a pressure of 30 MPa in Ar for the VC-doped   ZrB2-SiC powder using high-energy ball milling followed by hot-pressing [16]. Our recent study showed  that highly dense ZrB2-SiC composites with VC additives were obtained by hot-pressing at 1750°C and 20  MPa for 1 h in a vacuum [17]. The resulting composites  exhibited good mechanical properties [17]; however,  oxidation and its effect on strength are not well understood. Although oxidation of ZrB2-SiC composites has  been studied by many  researchers  [5-7,11,12],  the  oxidation behavior of the ZrB2-SiC composites with  VC additives should be unlike the ZrB2-SiC composites  without VC.  Therefore,  it  is necessary  for high temperature structural application to examine oxidation of the ZrB2-SiC composites with VC additives and  strength  retention after  the oxidation at elevated  temperature. The aim of this study was to examine the oxidation  and its effect on room-temperature flexural strength in  the hot-pressed ZrB2-SiC composites with VC additives  exposed  to air at different  temperatures between  1100ºC and 1300ºC, for up to 100 h.  2.  Experimental procedure  2.1. Composite materials  The composite materials used in this study were fabricated by mixing ZrB2 powder (d50 = 2.1 µm, Grade F,  Japan New Metals, Osaka) with 20 vol% α-SiC  = 0.5 µm, UF-15, H.C. Starck, Berlin, Germany) in a ball  mill with ethyl alcohol. In order to examine the effect of  the amount of added VC on oxidation of the composites,   (d50   CONTACT Shuqi Guo  305-0047, Japan  GUO.Shuqi@nims.go.jp   Research Center for Structural Materials, National Institute for Materials Science, Tsukuba, Ibaraki   © 2020 The Author(s). Published by Informa UK Limited, trading as Taylor & Francis Group on behalf of The Korean Ceramic Society and The Ceramic Society of Japan.  This is an Open Access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/4.0/), which permits  unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.  \\x0c', '2  S. GUO  3, 5, and 7 wt% VC (d50 = 1.8 µm, Japan New Metals) was  added to the ZrB2 and SiC powder mixture. The three  composition materials were fabricated by hot-pressing at  1750°C for 60 min under an applied pressure of 20 MPa in  a vacuum  from the ball-milled ZrB2 and SiC powder  mixtures with VC additives. The detailed hot-press procedure was  reported elsewhere  [17]. Relative densities  exceeding 98% were obtained for the three compositions  after the sintering [17]. Hereafter, the three composition  materials are denoted as ZSVC03, ZSVC05, and ZSVC07,  respectively.  2.3. Bending test  The flexural strengths of the as-sintered and the postoxidized specimens were measured at room temperature in a four-point bending test fixture (inner span  10 mm and outer span 20 mm). The bending test was  performed using an Autograph testing system (AG 50KNI, Shimadzu, Kyoto,  Japan) with a crosshead  speed of 0.5 mm/min. After testing, the fracture surface  of the specimen was examined using field emission  scanning electron microscopy (FE-SEM).  2.2. Oxidation test  Specimens, averaging 25 mm × 2.5 mm × 2 mm in size,  were cut from the sintered composite plates. The surfaces of the specimens were ground with an 800-grit  diamond wheel. The obtained specimens were ultrasonically cleaned in acetone and then kept in an oven  at a constant temperature of 100°C before oxidation.  Oxidation tests were performed at 1100°C, 1200°C, and  1300°C in air over a period of 100 h. The specimens  were supported on porous Al2O3 knife-edged fixtures  in an electronic furnace (SSFT-1520, Nikkato Co., Ltd.,  Tokyo, Japan). Before and after oxidation, the specimens were weighed using an analytical balance (MT5,  Mettler Toledo Co., Ltd., Switzerland) with an accuracy  of ± 0.001 mg. After oxidation, the specimen surfaces  were examined using X-ray diffraction (XRD) to identify  the oxidation products produced during the exposure  at elevated temperatures.  3. Results and discussion  3.1. Oxidation kinetics  The weight gain data from the oxidation of the three  compositions between 1100°C and 1300°C are shown  in Figure 1. The three compositions show a similar  oxidation behavior: the specific weight gain increases  rapidly within the initial state of oxidation and then the  specific weight gain increases gradually with oxidation  time. Plots of  the  square of  the weight gain as  a function of oxidation time for the three compositions  after oxidation at temperatures of 1100−1300°C are  shown in Figure 2. It is found that the oxidation of  the composites does not display parabolic kinetics at  the  initial stage of oxidation  (within ~10 h)  for all  temperatures. After ~10 h of oxidation, however, the  specific weight gain vs time curves approximate the  classical parabolic behavior. Similar oxidation behavior  was previously reported in Er2O3-doped SiC ceramic   Figure 1. Plots of specific weight gain as a function of time for the three compositions between 1100°C and 1300°C.  \\x0c', 'JOURNAL OF ASIAN CERAMIC SOCIETIES  3  Figure 2. Square of weight gain as a function of time for the three compositions between 1100°C and 1300°C.  oxidized between 1200°C and 1400°C in air [18] and  ZrB2-MoSi2-SiC composites oxidized at 1500°C in the  air [19]. When oxidation of a material at elevated temperature approximates to the classical parabolic behavior, the specific weight gain, W, vs oxidation time, t, at  a given temperature can be represented by the following equation [20,21]:   W 2 ¼ kt  (1)   where k is the parabolic rate constant which is dependent  on  the  temperature  and  follows  the  Arrhenius law:   �  k ¼ k0 exp  \\x00  �  Q  RT  (2)   where k0 is a frequency factor, Q is the oxidation activation energy (both k0 and Q are material constants), T is  the absolute temperature, and R is the gas constant. The  k values at different temperatures can be determined  from the slopes of the straight lines in Figure 2. Note that  the parabolic rate constant was determined for the period  of 10−100 h in which the parabolic oxidation behavior was  shown (Figure 2). The obtained k values for the three  compositions are summarized in Table 1. In addition, the  total specific weight gain data for the three compositions  after 100 h of oxidation at temperatures of 1100−1300°C  are also  included  in Table 1. At 1100°C, the k value   increased with the  increase of VC content. At 1200°C  and 1300°C; however,  the k values of ZSVC07 were  lower than those of ZSVC03, but they were higher than  those of ZSVC05. In addition, it is found that the k values  were greater at 1200°C and 1300°C  than at 1100°C  regardless of the VC content. In particular, the k values  at 1300°C were an order of magnitude greater than those  at 1100ºC. The plots of the k values as a function of  oxidation temperature are shown in Figure 3. The apparent oxidation activation energy was determined from the  data shown in this figure to be ~180 kJ/mol for ZSVC03,  ~120 kJ/mol for ZSVC05, and ~70 kJ/mol for ZSVC07  oxidized between 1100ºC and 1300ºC in air.  3.2. Oxidation products  Figure 4 shows XRD patterns of the specimen surfaces  for the three compositions before and after 100 h of  oxidation at temperatures of 1100−1300ºC. Before oxidation, the as-sintered composites comprised ZrB2, SiC,  V3B4, and ZrC phases for the three compositions, with  no VC. The presence of the V3B4 and ZrC phases is  a result of the reaction of VC with ZrB2 and ZrO2  impurity of the ZrB2 particle surface during the sintering [17]. The reaction led to the complete consumption  of the added VC during the sintering [17]; therefore,  the absence of VC in the resulting composites. Unlike   \\x0c', '4  S. GUO  Table 1. Parabolic rate constant and specific weight gain for the three compositions after 100 h oxidation between 1100°C and  1300°C.  Materials  Amount of VC (wt%)  ZSVC03 ZSVC05 ZSVC07  3 5 7  Parabolic rate constant, k (mg2 cm−4 h−1)  Total specific weight gain, W (mg cm−2)  1100°C  0.076 0.118 0.471  1200°C  1.805 0.777 1.321  1300°C  7.185 2.512 3.022  1100°C  2.8482 3.6402 7.0394  1200°C  13.437 8.881 11.813  1300°C  26.937 15.951 17.859  Figure 3. Arrhenius plots of oxidation rate constant as a function of temperature for the three compositions.  the as-sintered composites, for the samples oxidized at  1100ºC, the ZrB2, m-ZrO2, V2O5, and ZrSiO4 phases  were detected  in each  instance, with no SiC, V3B4,  and ZrC. It is evident that these phases formed during  the oxidation except the ZrB2 phase, as a result of  oxidation of ZrB2, SiC, V3B4, and ZrC. It is known that  SiC oxidized to form silica (SiO2) glassy phase as it was  exposed to air at high temperature. However, no SiO2  was  detected  in  the  postoxidized  samples.  The  absence of SiO2 seems to be associated with the formation of ZrSiO4 which is produced by the reaction of  ZrO2 with SiO2 during oxidation exposure. This reaction would substantially consume the SiO2 from the  oxidation of SiC; consequently, a little quantity of the  residual SiO2 could not be detected by XRD. For the  samples oxidized at 1200ºC and 1300ºC, on the other  hand, the m-ZrO2, SiO2 (αor β-cristobalite), ZrV2O7,  and ZrSiO4 phases were detected in the three compositions. However, neither V2O5 nor ZrB2 peaks were  detected in each instance. The disappearance of ZrB2  peaks after oxidation at 1200ºC and 1300ºC indicated   that the signals of ZrB2 observed on the oxidized surfaces at 1100ºC were from the bulk ceramic beneath  the oxidized layer because the oxidation reaction substantially evolved with the  increase of temperature  (Figure 1). In addition, although XRD analysis shows  the presence of the crystalline SiO2 for the postoxidized samples at 1200ºC and 1300ºC, an amorphous  silica hump was observed around 22 deg as well. The  presence of the amorphous silica hump implies that  the amorphous silica phase due to oxidation of SiC  would be partially crystallized into cristobalite phase,  resulting in the presence of a mixture of the crystalline  silica and amorphous silica phase  in  the oxidized  surfaces. It is known that ZrB2, ZrC, and SiC oxidize to form  ZrO2,   SiO2, and B2O3 when they are exposed to air at elevated  temperature  [5-7,11,12,19].  In  addition,  although not well known, it should be presumed that  V3B4 oxidizes to form V2O5 and B2O3 as it is exposed to  air at elevated temperature. Thus, when an ZrB2-SiC   \\x0c', 'JOURNAL OF ASIAN CERAMIC SOCIETIES  5  Figure 4. XRD patterns of the specimen surfaces for the three compositions before and after 100 h of oxidation between 1100°C  and 1300°C; (a) ZSVC03, (b) ZSVC05, and (c) ZSVC07.  composite with VC additive was exposed to air at  elevated  temperatures,  the  individual  constituents  ZrB2, SiC, ZrC, and V3B4 oxidized according to the  following reactions:   SiC sð  Þþ1:5O2 gð  Þ¼ SiO2  sð  ÞþCO gð  Þ  ZrB2  sð  Þþ2:5O2 gð  Þ¼ ZrO2  sð  ÞþB2O3  lð Þ  ZrC sð  Þþ1:5O2 gð  Þ¼ ZrO2  sð  ÞþCO gð  Þ  (3)   (4)   (5)   2V3B4  sð  Þþ13:5O2 gð  Þ¼ 3V2O5  lð Þþ4B2O3  lð Þ  (6)   Because the ZrO2, SiO2, and V2O5 produced coexisted  in the oxidized surfaces during oxidation, thus the ZrO2  reacts with the SiO2 or V2O5 to form ZrSiO4 or ZrV2O7  accompanying the evolution of oxidation as follows:   ZrO2  sð  ÞþSiO2  sð  Þ¼ ZrSiO4  sð  Þ  ZrO2  sð  ÞþV2O5  lð Þ¼ ZrV2O7  sð  Þ  (7)   (8)   \\x0c', '6  S. GUO  The formation of ZrSiO4 by the reaction of SiO2 with ZrO2  at elevated temperature is documented in the literature.  Rosen and Muan have reported that the change in Gibbs  free energy, ΔG, for the formation of ZrSiO4 by the reaction of SiO2 with ZrO2 was in the range - 9.96 kJ/mol to -  5.02 kJ/mol between 1180°C and 1366°C [22]. In addition,  Ellison and Navrotsky showed that the ΔG for the formation of ZrSiO4  from SiO2 and ZrO2  increased with an  increase of temperature between 298 K and 1900 K [23],  with ΔG<0. Based on the previous study of ΔG [23], the  ΔG value of the formation of ZrSiO4 from ZrO2 and SiO2  (Equation (7)) was determined to be in the range of -  15.86 kJ/mol to - 13.56 kJ/mol between 1100°C and 1300°  C. It is evident that the reaction (7) is thermodynamically  favorable at temperatures of 1100-1300°C. Furthermore,  Veytizou et al. [24] showed the formation of ZrSiO4 from  ZrO2 and amorphous SiO2 between 1250°C and 1400°C.  An earlier study of ZrB2-SiC composite showed that ZrO2  reacted with SiO2 to form ZrSiO4 during thermal cyclic  exposure at 1200ºC and 1400ºC in air [25]. Thereby, it  would be expected that the reaction (7) might be occurring for the  investigated material during the thermal  exposure at temperatures of 1100-1300ºC.  On the other hand, previous studies in V2O5-ZrO2  system showed the formation of ZrV2O7 by reaction of  ZrO2 with V2O5 (Equation (8)) at or below 750ºC [26,27].  The ZrV2O7  is  thermal unstable at  temperature of  >750ºC; however,  in  situ X-ray diffraction  revealed  that  it decomposes  into m-ZrO2  and  liquid V2O5  (Melting point: Tm = 690ºC)  [27]; the solid m-ZrO2  phase was dissolved  in the  liquid V2O5 phase. The  decomposition of ZrV2O7 above 750ºC implies that it  would not be expected for oxidation temperatures in  the range of 1100ºC to 1300ºC used in this study. On  the other hand, according to the V2O5− ZrO2 phase  diagram [28] (not shown), it is seen the presence of the  peritectic  reaction,  Liquid + ZrO2 → ZrV2O7,  at  a temperature of ~750ºC. The peritectic reaction occurring at ~750ºC would  lead to the formation of the  liquid-oxidized reactive layer on the oxidized surfaces  of  the  specimens  during  the  thermal  exposure  between 1100ºC and 1300ºC. The m-ZrO2 content in  the oxidized region increased accompanying continuous oxidation of ZrB2 and ZrC, and m-ZrO2 should  precipitate out of the liquid V2O5 as its concentration  is oversaturation. In addition, when the oxidized specimens were cooled down to room temperature from  the oxidation temperatures, the ZrV2O7 phase formed  below 750ºC due to the occurrence of the peritectic  reaction. As a result, the ZrV2O7 was detected in the  oxidized samples at 1200ºC and 1300ºC (Figure 4). One  exception is the absence of the ZrV2O7 phase for the  oxidized samples at 1100ºC (Figure 4). The absence of  ZrV2O7 suggests that the peritectic reaction was inhibited on cooling. This behavior seems to be associated  with a rapider cooling rate for the samples oxidized at   1100ºC than for those oxidized at 1200ºC and 1300ºC  because of the thinner oxidized layer for the former  than for the latter.  3.3.  Effect of oxidation on flexural strength  Figure 5 shows the effect of oxidation temperature on  room-temperature flexural strength of the three compositions.  It  is  found  that after 100 h of oxidation  between 1100°C and 1300°C, the flexural strength of  the composites varies depending on the amount of VC  and oxidation temperature. After 100 h of oxidation at  1100°C,  the  flexural  strength  of  the  composites  increased; the  increase of strength from that before  the oxidation is ~40% for ZSVC03, ~34% for ZSVC05,  and ~18% for ZSVC07. After 100 h of oxidation at 1200°  C, approximately 39% and 17% increases of strength  were obtained for ZSVC05 and ZSVC07, respectively.  Approximately 8%  loss of strength was observed for  ZSVC03, however. Unlike the postoxidized specimens  at 1100°C and 1200°C, after 100 h of oxidation at 1300°  C, the strength of the three compositions decreased; the  reduction of strength from that before the oxidation is  ~53% for ZSVC03, ~8% for ZSVC05, and ~5% for ZSVC07. The fracture surface of the postoxidized composites  was examined under FE-SEM observations, which exhibited that the fracture behavior was closely linked to the  oxidation temperature and the amount of VC. Figure 6  shows the FE-SEM micrographs of the fracture surfaces of  the three compositions after 100 h of oxidation at 1100°C.  The oxidized  layer  thickness was determined  to be  ~50 μm for ZSVC03, ~65 μm for ZSVC05, and ~105 μm  for ZSVC07. Additionally, a thin dense glassy scale was  observed on  the oxidized  surfaces  for ZSVC03 and  ZSVC05 (Figure 6(a,b)), with no delamination of the glassy  scale from the bulk bodies. Previous studies in ZrB2-SiC or  Si3N4 ceramics showed that the outermost thinner dense  glassy scale formed during the oxidation process led to  the increase of strength after oxidation by a crack blunting process [25,29]. In addition, our earlier study of ZrB2 SiC composites with VC showed that the fracture origin  for the as-sintered samples was located at the flaws in the  surfaces  [17]. Thus, similar crack blunting would be  expected  in  the materials  investigated  in  this study  because of  the  formation of  the outermost  thinner  dense oxide scale on the oxidized surfaces (Figure 6).  Under  higher-magnification  FE-SEM  observations,  although several cracks were observed in the oxide layer  for ZSVC05 (indicated by arrows  in Figure 6(c)), these  cracks have not been extended into the bulk body, with  no contribution to the degradation of strength. Therefore,  a substantial  increase  in strength was observed  for  ZSVC03 and ZSVC05. Unlike ZSVC03 and ZSVC05, for  ZSVC07 the local delamination of the oxide layer was  observed in the oxidized specimen at 1100°C for 100 h  (indicated by arrows  in Figure 6(d)). Although  the   \\x0c', 'JOURNAL OF ASIAN CERAMIC SOCIETIES  7  Figure 5. Effects of oxidation temperature on room-temperature flexural strength of the three compositions after 100 h of  oxidation between 1100°C and 1300°C.  Figure 6. FE-SEM micrographs of the fracture surfaces of the three compositions after 100 h of oxidation at 1100°C; (a) ZSVC03, (b,  c) ZSVC05, and (d) ZSVC07.  \\x0c', '8  S. GUO  separation of the entire oxide layer from the bulk body  was not observed, the local delamination created the new  defects inside the oxidized layer, thereby no significant  increase of strength was observed for ZSVC07 (Figure 5). Figure 7 shows the FE-SEM micrographs of the fracture surfaces of the three compositions after 100 h of  oxidation at 1200°C. The oxidized  layer thickness  is  ~150 μm  for ZSVC03, ~120 μm  for ZSVC05, and  ~140 μm  for ZSVC07. For ZSVC03, a damage zone  was observed beneath the outermost oxide layer for  the postoxidized specimen (Figure 7(a)). In addition,  a number of cracks was observed  in the outermost  oxide layer. These defects led to the degradation of  strength after oxidation (Figure 5). Unlike ZSVC03, the  outermost dense glassy scale was observed on the  oxidized surfaces for ZSVC05 and ZSVC07 (Figure 7  (b-d)). Although the cracks and/or the delamination  of the oxide layer were observed for both the compositions  (Figure 7(c,d)),  these cracks have not been  extended  into the bulk body (Figure 7(c)) and only  the local delamination occurred in the oxidized layer  without separation from the bulk body (Figure 7(d)).  The oxidation phenomenon observed for ZSVC05 and  ZSVC07 specimens oxidized at 1200°C  is essentially  similar to that observed for the specimens oxidized at  1100°C (Figure 6(b-d)). As a result, their strengths after  100 h of oxidation at 1100°C and 1200°C are nearly the   same, with a significant increase of strength compared  to the unoxidized specimens (Figure 5). Figure 8 shows the FE-SEM micrographs of the fracture  surfaces of the three compositions after 100 h of oxidation  at 1300°C. The oxidized reactive region was significantly  extended compared to the post-tested specimens at  1100ºC and 1200ºC. The oxidized  layer  thickness  is  ~520 μm  for ZSVC03, ~275 μm  for ZSVC05,  and  ~350 μm for ZSVC07.  In addition, for ZSVC03, pores-,  and cracks-containing oxide layer was observed for the  postoxidized specimen (Figure 8(a)). Obviously, the presence of pores and cracks led to the substantial degradation of strength after oxidation (Figure 5). For ZSVC05 and  ZSVC07, although the outermost dense glassy scale was  observed for the postoxidized specimens (Figure 8(b,c)),  a porous damage region existed beneath the glassy scale  for ZSVC05  (indicated by arrows  in Figure 8(b)) and  a longer crack extended from the outermost glassy scale  into the bulk body for ZSVC07 (Figure 8(c)). A previous  study in Si3N4 ceramics showed that the degradation of  strength after oxidation is attributed to the formation of  a damaged  region beneath  the  oxide  scale  [30].  Presumably, the decrease  in strength for ZSVC05 and  ZSVC07 after 100 h of oxidation at 1300ºC should be  attributed either to the formation of the damage zone  beneath the oxide scale or to the formation of the larger  new defects in the oxidized layer.  Figure 7. FE-SEM micrographs of the fracture surfaces of the three compositions after 100 h of oxidation at 1200°C; (a) ZSVC03, (b,  c) ZSVC05, and (d) ZSVC07.  \\x0c', 'JOURNAL OF ASIAN CERAMIC SOCIETIES  9  Figure 8. FE-SEM micrographs of the fracture surfaces of the three compositions after 100 h of oxidation at 1300°C; (a) ZSVC03, (b)  ZSVC05, and (c) ZSVC07.  4. Conclusions  The oxidation and its effect on room-temperature flexural strength were examined in the hot-pressed ZrB2  -20 vol% SiC composites with 3-7 wt.% VC additives  within the temperature range of 1100-1300ºC, for up   to 100 h. The major results obtained are exhibited  below: (1) The oxidation of the three compositions was found  to display parabolic behavior for the period of 10-100 h.   \\x0c', '10  S. GUO  The parabolic rate constant depended on the VC content  and  temperature. Among  the  three  compositions,  ZSVC05 had the minimum parabolic rate constant within  the temperature range of 1100-1300ºC. At 1100ºC, the  parabolic rate constant of ZSVC03 was smaller than that  of ZSVC07. At 1200ºC and 1300ºC, however, this constant  was greater for ZSVC03 than for ZSVC07. (2) The oxidation products of the three compositions  consisted  of m-ZrO2, V2O5,  and  ZrSiO4  at  1100ºC, m-ZrO2, ZrV2O7, ZrSiO4, and SiO2  (αor β cristobalite) at 1200ºC and 1300ºC. (3) The room-temperature strength of the three compositions increased after 100 h of oxidation at 1100ºC.  After 100 h of oxidation at 1200ºC, ZSVC05, and ZSVC07  retained  the  increased  strength, but  the  strength  decreased  for ZSVC03. After 100 h of oxidation at  1300ºC, however, the strength degraded for the three  compositions; the substantial degradation of strength  was observed only for ZSVC03.  Disclosure statement  The authors declare that they have no conflict of interest.  References  [1] Fahrenholtz WG, Hilmas GE, Talmy IG, et al. Refractory  diborides of zirconium and hafnium. J Am Ceram Soc.  2007;90:1347-1364.  [2] Wuchina E, Opila E, Opeka M, et al. UHTCs: ultra-high  temperature ceramic materials for extreme environment applications. Interface. 2007;16:30-36.  [3]  [5]  [7]  Paul A, Jayaseelan DD, Venugopal S, et al. UHTS composites for hypersonic applications. Am Ceram Soc  Bull. 2012;91:22-29. [4] Matkovich VI. Boron and refractory borides, metallic  borides: preparation of solid bodies, sintering methods and properties of solid bodies. New York (NY):  Springer; 1977. p. 457-493. Kuriakose AK, Margrave JL. Oxidation kinetics of zirconium  diboride  and  zirconium  carbide  at  high  temperatures. 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Effects of VC additives on densification and  elastic and mechanical properties of hot-pressed ZrB2  -SiC composites. J Mater Sci. 2018;53:4010-4021. [18] Guo SQ, Hirosaki N, Tanaka H, et al. Oxidation behavior  of liquid-phase-sintered SiC with AlN and Er2O3 additives between 1200ºC and 1400ºC. J Eur Ceram Soc.  2003;23:2023-2029. [19] Guo SQ, Mizuguchi T,  Ikegami M, et al. Oxidation  behavior of ZrB2-MoSi2-SiC  composites  in  air  at  1500ºC. Ceram Int. 2011;37:585-591. [20] Singhal SC. Thermodynamics and kinetics of oxidation of  hot-pressed silicon nitride. J Mater Sci. 1976;11:500-509. [21] Cinibulk MK, Thomas G, Johnson SM. Oxidation behavior of  rare-earth disilicate-silicon nitride ceramics.  J Am Ceram Soc. 1992;75:2044-2049. [22] Rosen E, Muan A. Stability of zircon in the temperature  range  1180°C  to  1366°C.  J  Am  Ceram  Soc.  1965;48:603-604. [23] Ellison AJG, Navrotsky A. Enthalpy of  formation of  zircon. J Am Ceram Soc. 1992;75:1430-1433. [24] Veytizou C, Quinson JF, Valfort O, et al. Zircon formation  from amorphous silica and tetragonal zirconia: kinetic  study and modeling. Solid State Ion. 2001;139:315-323. [25] Guo SQ. Strength retention in hot-pressed ZrB2-SiC composite after thermal cycling exposure in air at 1200ºC and  1400ºC. J Am Ceram Soc. 2019;102:3842-3848. [26] Buchanan RC, Wolter GW. Properties of hot-pressed  zirconium pyrovanadate ceramics. J Electrochem Soc.  1983;130:1905-1910. [27] Chen Z, Speakman S, Howe J, et al. Investigation of reactions between vanadium oxide and plasma-sprayed  yttria-stabilized  zirconia coatings.  J Eur Ceram Soc.  2009;29:1403-1411. [28] Ondik HM, McMurdie HF, editors. Phase diagrams for  zirconium + zirconia systems. Westerville (OH): The  American Ceramic Society; 1998. p. Fig. Zr-212. [29] Jakus K, Ritter JE, Rogers WP. Strength of hot-pressed  silicon nitride after high temperature exposure. J Am  Ceram Soc. 1984;67:471-475. Park H, Kim HW, Kim HE. Oxidation and strength retention of monolithic Si3N4 and nanocomposite Si3N4-SiC  with Yb2O3 as a sintering aid.  J Am Ceram Soc.  1998;81:2130-2134.  [30]  \\x0c']"
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  "PDF": "Oxidation and strength retention of HfB2−SiC composite with La2O3 additives.pdf",
  "Text": "['Advances in Applied Ceramics  Structural, Functional and Bioceramics  ISSN: 1743-6753 (Print) 1743-6761 (Online) Journal homepage: https://www.tandfonline.com/loi/yaac20  Oxidation and strength retention of HfB2−SiC composite with La2O3 additives  Shuqi Guo  To cite this article: Shuqi Guo (2020): Oxidation and strength retention of HfB2−SiC composite with La2O3 additives, Advances in Applied Ceramics, DOI: 10.1080/17436753.2020.1755510  To link to this article:  https://doi.org/10.1080/17436753.2020.1755510  Published online: 23 Apr 2020.  Submit your article to this journal   Article views: 3  View related articles   View Crossmark data  Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=yaac20  \\x0c', 'Oxidation and strength retention of HfB2−SiC composite with La2O3 additives  Shuqi Guo  Research Center for Structural Materials, National  Institute for Materials Science, Tsukuba, Japan  ABSTRACT  The oxidation and its effect on room temperature flexural strength were examined in hotpressed 3 vol.-% La2O3-doped HfB2-20 vol.-% SiC composite exposed to air at 1400°C and 1500°C, respectively, for up to 100 h. The major oxidation products consisted of SiO2 (αcristobalite), La2Si2O7, HfO2 and HfSiO4 regardless of oxidation temperature. However, the morphology of the microstructure of the oxidised surfaces was shown to be dependent on oxidation temperature. In addition, it was found that the room temperature flexural strength of the composite degraded after oxidation and the degradation of strength depended on oxidation time at temperature and oxidation temperature. The strength retention displayed by composite after 100 h of oxidation at 1400°C and 1500°C was 83% and 74%, respectively.  ARTICLE HISTORY  Received 30 January 2020 Revised 17 March 2020 Accepted 7 April 2020  KEYWORDS  Hafnium diboride; Silicon carbide; composite; oxidation; flexural strength  Introduction  Zirconium diboride  (ZrB2) the most  and  hafnium diboride  (HfB2) are both of family of ultra-high temperature  important members of a  ceramics  (UHTCs).  They have extremely high melting points  (>3000°C),  high thermal conductivities, great strength and ablation resistance [1-4]; these features make them become  an important class of UHTCs for a variety of thermo mechanical structural applications at elevated tempera tures. However,  ZrB2 and HfB2 insufficient oxidation resistance for hypersonic flight  ceramics  possess  environments,  as  they  start  oxidising  at  or  greater  800°C [5,6]. One  of  the  promising  approaches  for  improving their oxidation resistance is to incorporate  SiC into ZrB2 and HfB2 ceramics to form a protective glassy borosilicate layer in oxidising environments [7-11].  In  addition,  the  addition  of  SiC particles  improved  room and  high-temperature ZrB2 and HfB2 ceramics [11-16]. Thus, the presence of SiC particles in the ZrB2- and HfB2-SiC composites is indispensable to obtain excellent mechanical proper strengths  of  ties,  better  oxidation, resistance [7-16]. Recently, HfB2-SiC composites potential candidates for the thermal protection systems  thermal  shock  and ablation ZrB2- attractive  both  the  and  are  becoming  and  of re-entry hypersonic vehicles with sharp leading edge In particular, HfB2-SiC composites have higher resistance to ablation/oxidation than ZrB2-SiC composites [7-10,18,19]. Furthermore, higher density of HfB2-SiC composites centre of mass of hypersonic vehicles to be changed,  profiles  [17,18].  the  enables  the  thereby  enhancing  the manoeuvrability  during reHfB2-SiC  entry  [20].  These  advantages  make  composites  a more  attractive  alternative  than ZrB2-  SiC composites.  Unfortunately, at a temperature equal  to or greater  2000°C, the glassy borosilicate layer is lost in a very short period (for 1 h at 2000°C in a non-air flowing atmosphere) and subsequently the underlying bulk  material is exposed to continuous oxidation, as a result at 1650°C. SiO2 improve oxidation resistance, it is necessary to raise  of melting  of  pure  To  further  the refractoriness and/or the melting point of the pro tective glassy borosilicate layer. Previous study in ZrB2-SiC composites with 10 wt-% La2O3 or Gd2O3 showed that oxidation for 1 h at 1600°C in static air  led  to  the  formation  of  an  outermost  dense ZrO2 layer and intermediate layers of heterogeneous RE2Si2O7 (RE: Rare Earth) and amorphous silicate oxides, instead of a porous ZrO2 layer for the ZrB2-SiC composite without RE-additives [9]. A latter study in  ZrB2and HfB2-20 vol.-% SiC composites with 2 wt% La2O3 additions showed that La2O3 addition increased glassy borosilicate melt viscosity and lowered  the oxygen diffusion coefficient through the glassy bor osilicate melt which was  formed in the oxidised sur faces of  the composites exposed to air between 1400°  C and 1600°C [21],  therefore higher oxidation resist ance. However,  the  retained fracture  strength of  the  composites after oxidation for  longer period at high  temperatures in air is not well-known. Strength reten tion after oxidation is an important high-temperature  property, because these materials currently are being  considered for use as structural components  in high temperature oxidising environments. The aim of  the  present  study was  to examine the effect of oxidation  © 2020 Institute of Materials, Minerals and Mining. Published by Taylor & Francis on behalf of  the Institute.  CONTACT Shuqi Guo GUO.Shuqi@nims.go.jp kuba, Ibaraki 305-0047, Japan  Research Center for Structural Materials, National  Institute for Materials Science, 1-2-1 Sengen, Tsu ADVANCES IN APPLIED CERAMICS https://doi.org/10.1080/17436753.2020.1755510  \\x0c', 'on the room temperature strength of HfB2-SiC composite with La2O3 additive.  hot-pressed  Experimental procedure  The starting powders used in this (d50 = 4.1 µm, Grade O, α-SiC  study were: HfB2 Japan New Metals, Osaka, (d50 = 0.5 µm, UF-15, H.C. Starck, Berlin, Germany), La2O3 powder (99.9% pure, Kojundo Chemical Laboratory, Saitama, Japan). The  Japan),  powder  3 vol.-% La2O3-doped HfB2-20 vol.-% SiC composite was hot-pressed at 2000°C for 60 min under a uniaxial  load of 20 MPa in a flowing Ar atmosphere from ball milled mixture of HfB2, SiC and La2O3 powders. The theoretical densities of the composite were calculated  according to the  rule of mixtures. Density values of  11.21 g cm  −3  for HfB2, for La2O3 are used in the calculation of the theoretical density of the composite. The density  3.22 g cm  −3  for  SiC  and  6.51 g cm  −3  of  the as-sintered composite compact was determined  using the Archimedes method to be 9.43 g cm  −3;  the  relative density was approximately 99.6%.  Test  specimens  averaging 25 × 2.5 × 2 mm in size  were cut from the composite plates. All the specimens  were ground with an 800-grit diamond wheel. The specimen’s  surfaces were  ultrasonically  cleaned  in  acetone and kept in an oven with a constant tempera ture of 100°C before oxidation test. Oxidation tests  were performed at 1400°C and 1500°C in air over a  period of 100 h. The specimens were supported on por ous Al2O3 knife-edged fixtures in an electronic furnace (SSFT-1520, Nikkato, Tokyo, Japan). Before and after  oxidation,  the  specimens  were  weighed  using  an  analytical  balance  (MT5, Mettler Toledo Co.,  Ltd.,  Switzerland) with  an  accuracy  of ±0.001 mg. After  100 h of oxidation,  the specimens were examined by  X-ray diffraction (XRD)  to determine  the oxidation  products produced during the oxidation.  In addition,  the oxidised surfaces were characterised by field-emis sion scanning electron microscopy (FE-SEM).  The fracture strength of  the pristine and post-oxi dised specimens was determined by fracture at  room  temperature,  using  four-point  flexure  (inner  span  10 mm, outer  span 20 mm). The bend test was per formed with an Autograph testing system (Autograph  Model AG-50KNI,  Shimadzu, Kyoto, −1. At mens were used for each measurement. After the bend Japan) with a  crosshead speed of 0.5 mm min  least five speci ing test,  the fracture surface was examined using FE SEM.  Results and discussion  Figure 1 shows plots of the specific weight gains of the  specimens oxidised at 1400°C and 1500°C. As  is evi dent  in this figure,  it  is found that  the specific weight  gain in specimens  increases  rapidly within an initial  10 h of with further  oxidation and  then it  increases  gradually  increasing oxidation time for  the speci mens oxidised at 1400°C and 1500°C, regardless of oxi dation  temperature. A similar  time  dependency  of  weight gains implies that oxidation behaviour occurred  during the  exposure at 1400°C and 1500°C in air  is  essentially the same. This time dependency of weight  gains suggests that after  the outermost oxidised layer  formed at  the initial stage of oxidation,  the diffusion,  such as the outward diffusion of constituent elements  cations from the bulk to the oxidised surface and the  inward diffusion of O through the scale layer,  is  the  rate-controlling mechanism for oxidation. Addition ally,  the weight gain was significantly lower at 1400°C  than at 1500°C. After 100 h of oxidation at 1400°C  and 1500°C, the weight gain displayed by the composite was 4.19 and 6.11 mg cm −2, respectively. Figure 2 shows XRD patterns of the specimen’s sur faces before and after 100 h of oxidation at 1400°C and  1500°C. Before oxidation exposure (Figure 2(a)), HfB2, SiC and La2Si2O7 phases were clearly detected in the specimen’s surface, with no La2O3. HfB2 is the primary crystalline phase, SiC and La2Si2O7 are the secondary phase. The presence of La2Si2O7 phase is attributed to the reaction of SiO2 on the surfaces of SiC particles with La2O3 during the sintering. As a result, the La2O3 peak is the absence in the XRD curve of the as-sintered  composite (Figure 2(a)). For the post-oxidised samples  at 1400°C and 1500°C (Figure 2(b,c)), on the other hand, new peaks of HfO2, HfSiO4 and SiO2 (α-cristobalite) phases were detected in addition to La2Si2O7, with no HfB2 and SiC phases. Obviously, the presence of HfSiO4, HfO2 and SiO2 phases on the oxidised surfaces of HfB2-SiC-La2O3 material is associated with oxidation of HfB2 and SiC occurred during exposure at 1400°C and 1500°C in air. In addition, the intensity  of HfSiO4 peaks was substantially stronger at 1500°C than at 1400°C. This stronger intensity implies that  the increase of  temperature promoted the formation  Figure 1. Plots of specific weight gain as a function of time at 1400°C and 1500°C.  2  S. GUO  \\x0c', 'of HfSiO4 accompanying the acceleration of oxidation of HfB2 and SiC. Furthermore, the intensity of La2Si2O7 peaks was substantially stronger after oxidation than before oxidation. The stronger intensity indicated  that  La2Si2O7 formed due to oxidation compared to the as-sintered  phase  is  enriched  in the  oxide  scale  sample. However, no B-containing oxidation product  was detected for the oxidised surfaces of the specimens  at 1400°C and 1500°C for 100 h. Our earlier study in ZrB2-SiC-MoSi2 materials showed that no B-containing oxidation product was formed on the oxidised sur faces of specimens after oxidation at 1500°C in air [22].  The absence of B-containing oxidation product  indi cates  that  the B2O3 formed by the oxidation of HfB2 is completely vaporised at 1400°C and 1500°C because  B2O3 has a high vapour pressure and is vaporised above 1300°C [5,7,8].  Figure 3 shows the backscattered electron FE-SEM  morphologies of unoxidised and oxidised surfaces of  the  specimens at 1400°C and 1500°C for 100 h. For  the unoxidised specimen (Figure 3(a)),  it  is seen that  the  composite  consisted of  the  equiaxed HfB2 contrast), SiC (dark contrast) and the La2Si2O7 grainboundary phase (dark-grey contrast). For the post-oxi (grey  dised specimen at 1400°C (Figure 3(b)), on the other  hand, the oxidised surface contained the rare-earth dis ilicates (La2Si2O7) in the equiaxed and the needleand/ or platelike morphologies (white-grey contrast), finer  HfO2 and/or HfSiO4 the black background of SiO2 glassy phase. Unlike the oxidised surface  particles  (white contrast) (α-cristobalite)  and  and a  at 1400°C,  for  the post-oxidised specimen at 1500°C (Figure 3  (c)), the oxidised surface consisted almost of the disili(white-grey  cates  in the platelike morphology  con trast),  larger  equiaxed HfO2 and/or HfSiO4 (white contrast) and the black background of (α-cristobalite) and a glassy phase.  crystals  SiO2  It  is well-known that  the rare-earth disilicates were  formed in Si3N4 with rare-earth oxide additions when it was exposed to air at elevated temperatures. Lange  et al. [23] Si3N4-Y2O3 Y2Si2O7 in the oxide scale, as a result of diffusion of Y cation from the grain boundaries  studied  the  oxidation  of  a  hot-pressed  ceramic  and  observed  enrichment  of  the outward  in  Si3N4 Sanders [24] showed the formation of needle-like La2Si2O7 grain in the surface of Si3N4-La2O3 specimens oxidised at 1370°C in air for 200 h via diffusion of La  substrate to the oxide surface. Mieskowski and  cation from Si3N4 grain boundaries to the oxide scale during oxidation. In addition, Babini et al. [25] had  reported that  formation of Y2Si2O7 phase with the different morphologies during oxidation via dissol ution and precipitation of Y2Si2O7 glassy phase for Si3N4-Y2O3-SiO2 to air between 1000°C and 1400°C for up to 30 h.  in the  silica-rich  ceramics  exposed  They  concluded  that  the morphology  of  Y2Si2O7 compositional  depended  on  oxidation  temperature,  parameters and quenching rate. In general, at low oxi dation temperature, only high aspect ratio needle-like  crystals develop, whereas  the platelike  forms are  the  most  preferred  at  higher  oxidation  temperatures.  Additionally,  they found that a low concentration of  Figure 2. XRD patterns of  the specimen surface before and after 100 h of oxidation at 1400°C and 1500°C.  ADVANCES IN APPLIED CERAMICS  3  \\x0c', '4  S. GUO  Figure 4. Effect of oxidation time on room temperature flexural strength after oxidation at 1400°C and 1500°C.  La cations  from the bulk to the glassy scale and the  inward diffusion of O through the  scale  layer. The  different morphology of La2Si2O7 observed at dised surfaces at 1400°C and 1500°C is associated  the oxi with the  concentration of  La  cation and  oxidation  temperature. On the other hand, earlier studies in ZrB2-SiC and ZrB2-SiC-La2O3 composites showed the formation of the ZrSiO4 phase after oxidation at high temperatures in air. Chen et al. have reported the presence of post-oxidised ZrB2-SiC-La2O3 ZrSiO4 composite at 1500°C and 1600°C in air [26]. Our study of ZrB2-SiC composite of ZrSiO4 phase during the thermal cyclic exposure at 1200°C and 1400°C in air as well [27]. The formation  showed the  formation  in the  phase  is  phase  reaction of  associated with the  of ZrSiO4 ZrO2 with SiO2 during oxidation at elevated temperatures in air [26,27]. Similar reaction between HfO2 and SiO2, although not well-known, could be expected during oxidation at 1400°C and 1500°C for the HfB2- SiC-La2O3 material investigated in this study. However, a previous study in HfB2-20 vol.-% SiC-2 wt-% La2O3 composite consolidated by spark plasma sintering showed that HfSiO4 phase was absent in the oxidised samples at 1400°C and 1500°C in air for ≤32 h [21]. Obviously, a further detailed investigation of the  formation of HfSiO4 phase during oxidation at high temperatures in air is clearly needed.  Figure 4 shows the effect of oxidation on the room  temperature flexural  strength of  the composite. As  is  evident  in this figure,  the strength retention displayed  by  the  composite  after oxidation depended on oxi dation time  and temperature. For  the post-oxidised  specimens  at  1400°C,  the  flexural  strength  initially  increased and then decreased with increasing oxidation  time. The flexural  strength increased from 580 MPa  before oxidation to 630 MPa after 10 h of oxidation at 1400°C for increase of strength 10%. Subsequently, the flexural strength decreased to 510 and 480 MPa  Figure 3. Backscattered electron FE-SEM micrographs of the specimen’s surfaces (a) before and after 100 h of oxidation at (b) 1400°C and (c) 1500°C.  Y3+ in the surface silicate melts and slow mobility of Y3 + mobility  the needle-like  favourable  crystal  are  to  form.  In the present  concentration of La cation in the glassy scale is lower at 1400°C (La3+: 0.05 at.-%) than at 1500°C (La3+: 0.24 at.-%). Presum study,  the  ably,  the formation of La2Si2O7 in the oxide scale is a result of the dissolution and precipitation of La2Si2O7 in the silica-rich glassy scale during oxidation at 1400°C and 1500°C for the HfB2-SiC-La2O3 material investigated in this study, by the outward diffusion of  \\x0c', 'after 50 and 100 h of oxidation respectively for strength retentions 88% and 83%. specimens at 1500°C, on the other hand,  For  the  post-oxidised  the flexural  strengths decreased with increasing oxidation time at  temperature, with no increase of strength. The strength  reduced from 580 MPa before oxidation to 530, 480  and  430 MPa  after 10, 50 and 100 h of oxidation strength retentions 90%, 82% and  respectively for 74%. The fracture surfaces of the post-oxidised specimens  at 1400°C and 1500°C were observed by FE-SEM; typi cal fracture appearances are shown in Figure 5. For the  oxidised specimen at 1400°C for 10 h,  a  thin dense  glassy scale was observed on the oxidised surface of  the  specimen (Figure 5(a)), with no visible pores  in  the oxidised layer and/or delamination of  the outer most glassy scale.  It  is evident  that  the formation of  the thin dense glassy scale can heal  the surface flaws  without creating new cracks and/or defects at  the oxi dised surface,  therefore  increase of  strength (Figure  4). For the post-tested specimen at 1400°C for 100 h,  however,  the  outermost  glassy  scale  delaminated  from the thick oxidised layer (indicated by arrows in  Figure 5(b)). The delamination behaviour of the outer most  glassy  scale was  observed  for  the  post-tested  specimens oxidised at 1500°C as well, regardless of oxi dation  time  (indicated  by  arrows  in  Figure  5(c,d).  Additionally,  for the oxidised specimen at 1500°C for  100 h, several  large pores were clearly observed in the  outermost glassy scale (indicated by arrows in Figure  5(d)). The degradation of strength due to exposure to  air  at  elevated temperatures  is well documented in  the literature. Losses in flexural strength of more than  20% and  30% have  been  reported  in  Si3N4 with after 100 h of oxidation  Yb2O3 exposure at 1400° and 1500°C in air [28]. Our recent study in ZrB2-SiC composite showed that strength reduced by 20% and 55%, respectively, after cyclic exposure at 1200°C for 1000 cycles  additive  respectively  the flexural  and  1400°C for 500 cycles in air [27]. The loss in strength  after oxidation exposure  at  elevated temperatures  is  attributed to the presence of  either  larger pores or  defects  in the oxidised layer and the delamination of  the  outermost  thinner  dense  glassy  scale  [27,28].  Thus,  it  is reasonable to expect  that  the delamination  of  the outermost  glassy  scale  and the  formation of  pores  in the glassy layer may be the major cause of  strength  decrease  after  oxidation  for  the  studied  material.  Summary  The oxidation products of HfB2-SiC composite with La2O3 additive consisted of SiO2, La2Si2O7, HfO2 and HfSiO4 after 100 h of oxidation at 1400°C and 1500° C. The strength retention of the composite after oxi dation at 1400°C and 1500°C varied with oxidation  temperature and time at temperature. For the post-oxi dised  specimens  at  1400°C,  the  room temperature  flexural  strength  initially  significantly  increased,  Figure 5. FE-SEM micrographs of and (d) 1500°C, 100 h.  fracture surfaces of the oxidised specimens;  (a) 1400°C, 10 h,  (b) 1400°C, 100 h,  (c) 1500°C, 10 h  ADVANCES IN APPLIED CERAMICS  5  \\x0c', '6  S. GUO  subsequently  the  strength  reduced with  increasing  exposure  time.  For  the  post-oxidised  specimens  at  1500°C,  however,  the  flexural  strength  significantly  degenerated with increasing exposure time at tempera ture. The strength retention in the specimens oxidised at 1400°C and 1500°C for 100 h was 83% and 74%, respectively. The strength degradation was attributed  to the  formation of pores  in the oxidised layer  and  the delamination of the outermost thinner dense glassy  scale.  Disclosure statement  conflict  of  interest  was  reported  by  the  No potential author(s).  References  J  [6]  [5]  [4]  [1]  al. J  for Am  IG, et hafnium.  Fahrenholtz WG, Hilmas GE, Talmy Refractory diborides of zirconium and Am Ceram Soc. 2007;90:1347-1364. [2] Kalish D, Clougherty EV, Kreder K. Strength, fracture mode, and thermal stress resistance of HfB2 and ZrB2. J Am Ceram Soc. 1969;52:30-36. [3] Upadhya K, Yang JM, Hoffmann WP. Materials ultrahigh temperature structural applications. Ceram Soc Bull. 1997;76:51-56. 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Compressive creep and oxidation resistance of an Si3N4 material fabricated in the system Si3N4-Si2N2O-Y2Si2O7. J Am Ceram Soc. 1983;66:C98-C99. [24] Mieskowski DM, Sanders WA. Oxidation of silicon nitride sintered with rare-earth oxide additions. J Am Ceram Soc. 1985;68:C160-C163. Babini GN, Bellosi A, Vincenzini P. Factors influencing structural evolution in the oxide of hot-pressed Si3N4Y2O3-SiO2 materials. J Mater Sci. 1984;19:3487-3497. [26] Chen M, Li H, Yao X, et al. High temperature oxidation resistance of La2O3-modified ZrB2-SiC coating for SiCcoated carbon/carbon composites. J Alloys Compd. 2018;765:37-45. [27] Guo SQ. Strength retention in hot-pressed ZrB2-SiC composite after thermal cycling exposure in air at 1200°C and 1400°C. J Am Ceram Soc. 2019;102:3843-3848. Park H, Kim HW, Kim HE. Oxidation and strength retention of monolithic Si3N4 and nanocomposite Si3N4-SiC with Yb2O3 as a sintering aid. J Am Ceram Soc. 1998;81:2130-2134.  [21]  [23]  [25]  [28]  \\x0c']"
},{
  "_id": 137,
  "PDF": "Oxidation behavior and ablation mechanism of Cf-ZrB2-SiC composite fabricated by vibration-assisted slurry impregnation combined with low-temperature hot pressing.pdf",
  "Text": "['Corrosion Science 161 (2019) 108181  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Oxidation behavior and ablation mechanism of Cf/ZrB2-SiC composite fabricated by vibration-assisted slurry impregnation combined with lowtemperature hot pressing  T  Dongyang Zhang, Ping Hu⁎, Shun Dong, Xu Liu, Chenglin Wang, Zenan Zhang, Xinghong Zhang⁎  National Key Laboratory of Science and Technology on Advanced Composites in Special Environments, Harbin Institute of Technology, Harbin, 150001, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Cf/ZrB2-SiC Thermal shock resistance Oxidation Protective layer  1.  Introduction  The Cf/ZrB2-SiC composite was fabricated by vibration-assisted slurry impregnation and low-temperature hot pressing. The Cf/ZrB2-SiC composite exhibited an excellent thermal shock resistance and oxidation resistance with a dense SiO2-rich glassy layer exposing at 1500 °C/1 h condition for static oxidation. Moreover, the high temperature ablation resistance was investigated with an oxyacetylene ablative testing at temperature 1800 °C for 1000s. The mass loss rate and line loss rate were 7.32×10−4 mg/mm2·s and -1.70×10−4 mm/s, respectively. The enriched silica glass layer embedded with a small amount of zirconia particles was contributed mainly to the excellent ablation resistance of Cf/ZrB2-SiC composite.  The impelling demand for higher mach numbers of hypersonic reentry vehicles and propulsion applications, require thermal protective materials for such applications could withstand extreme temperature (> 2000 °C) in corrosive/oxidative atmospheres with high thermal gradient and achieve sufficient mechanical resistance [1-3]. The ultrahigh temperature ceramics (UHTCs), such as ZrB2, HfB2, ZrC and HfC have been identified as potential candidates owing to their high melting [4-6]. ZrB2 point and excellent oxidation ablation resistance as an important member in UHTC family contains many outstanding properties such as high melting point, high hardness and excellent mechanical property in high temperature [7,8]. The addition of SiC phase could enhance the densification via inhabiting grain growth and improve oxidation resistance by the oxidized dense SiO2 layer [9,10]. However, the low fracture toughness and poor thermal shock resistance prevent its engineering application. The addition of carbon fiber, especially continuous carbon fiber, could be greatly enhance the thermal shock resistance and improve the damage tolerance via multiple toughening mechanisms (fiber pull-out, fiber bridging, crack deflection) [11-13]. For the continuous carbon fiber reinforced ceramics, the UHTCs phases, such as ZrB2, HfB2, ZrC or HfC was usually introduced into carbon fiber preform by slurry impregnation (SI) [14] and precursor infiltration and pyrolysis (PIP) [15]. However, the relative density was limited and the closed porosity was  approached to 10% or higher [16]. The reactive metal infiltration process was en effective method to fabricate carbon fiber reinforced ultra-high temperature ceramic composites (Cf/UHTCs) with low open porosity, however, the reacted temperature was approached to 1500 °C or higher, leading to the degradation of carbon fiber [17,18]. Moreover, the residual low melting phases even in low amount was detrimental for the mechanical properties. Fortunately, Sciti group developed a hybrid process including typical vacuum-bag infiltration and conventional hot pressing to fabricate continuous carbon fiber reinforced ceramic composite [19,20]. The continuous carbon fiber reinforced ZrB2-SiC composite exhibited a laminated structure, showing a novel unidirectional properties in thermal shock resistance and mechanical properties [21,22]. Moreover, those composites were based on ultra-high temperature ceramics baseline and exhibited a superior ablation oxidation resistance [23,24]. However, the carbon fiber reinforcement in the Cf/ UHTC composite was fiber cloths or fiber bundles, exhibited a laminated structure with unidirectional properties. The introduction of three-dimensional carbon fiber preform into ceramic matrix could improve the unidirectional mechanical properties of Cf/UHTC composites and exhibited an excellent comprehensive properties. In our previous work, we reported a novel method to fabricate Cf/ZrB2-SiC composites by vibration-assisted slurry impregnation and low-temperature (1450 °C) hot pressing, which the mechanical properties and work of fracture were dramatically enhanced [14]. However, the investigation on thermal shock resistance and oxidation resistance of Cf/ZrB2-SiC  ⁎ Corresponding authors. E-mail addresses: huping123hit@163.com (P. Hu), zhangxh@hit.edu.cn (X. Zhang).  https://doi.org/10.1016/j.corsci.2019.108181 Received 2 June 2019; Received in revised form 22 August 2019; Accepted 25 August 2019  Available online 27 August 2019 0010-938X/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  composite were limited. In the present work, the Cf/ZrB2-SiC composite was achieved by the vibration-assisted slurry impregnation and low-temperature hot pressing. The microstructures and thermal shock resistance of the Cf/ ZrB2-SiC composite were analyzed and discussed in detail. The ability of ZrB2-SiC system to protect carbon fiber from static oxidation and dynamic ablation condition was analyzed in detail. The possible ablation mechanism was also proposed based on the characterization results of the scanning electron microscopy (SEM), energy-dispersive spectroscopy (EDS) and X-ray diffraction (XRD).  2. Experimental procedures  2.1. Materials preparation  The detailed fabrication process of the Cf/ZrB2-SiC composite was shown in our previous study [14]. Firstly, the carbon fiber preform with a density of 0.15 g/cm3 was fabricated by a fine weaving process, then the mixed ZrB2-SiC slurries were filled into the fiber preform by vibration-assisted slurry impregnation. The dried Cf/ZrB2-SiC green body was obtained after vacuum drying at 25 °C for 2 h, and the specific fiber volumetric amount was approached to 29 vol.%. Finally, the high-performance Cf/ZrB2-SiC composite was achieved successfully after lowtemperature hot pressing (1450 °C/30 MPa/2 h).  2.2. Thermal shock resistance  The thermal shock resistance of the Cf/ZrB2-SiC composite was investigated by a water-quenching technique. The samples were heated up to the thermal shock temperature from 125 °C to 925 °C in a muffle furnace and maintained at the specified heating temperature for 15 min to eliminate the effect of temperature gradient. The heated specimens were then subjected to a thermal shock by quenching them into a water bath, and the samples after thermal shock testing was used for residual strength testing. The residual flexural strength of thermal shock samples was measured by the three-point bending tests (Model 5569, Instron, USA) using the International Standards of ISO 14,704: 2016, and the size of testing sample was 3 mm × 4 mm × 36 mm with a span of 30 mm and a cross-head speed of 0.5 mm/min at room temperature. (ΔTc) The measured critical thermal shock temperature difference value was defined as 70% of the original flexural strength at room temperature according to the linear interpolation of the residual strength values, as described in ASTM C1525-04 [25]. The finite element analysis (FEA) was used to simulate the temperature and thermal stress distribution in the samples during water quenching process in different temperature. The thermal diffusivity, specific heat and thermal conductivity of Cf/ZrB2-SiC in room temperature were 3.68 mm2/s, 0.53 J/g·K and 7.73 W/m·K, respectively. The heat transfer coefficient of 130 kW/(m K) was used for the water quenching calculations.  2.3. Ablation testing  The oxidized Cf/ZrB2-SiC sample with dimension 15 mm × 10 mm × 2.5 mm (Length × Width × Thickness) was machined from the sintered pellet, and the oxidation resistance of the sample was evaluated by static oxidation at 1500 °C for 1 h in a muffle furnace. Moreover, the oxyacetylene ablative testing was performed to analyze the anti-ablation property of Cf/ZrB2-SiC composite according to GJB323A-96 [26]. The flux of C2H2 and O2 were 0.6 m3/h and 0.6 m3/h, respectively. The sample size for the ablative testing was Φ30 mm × 10 mm and the sample was placed vertically to the oxyacetylene torch with a distance of 30 mm, and the ablative temperature and ablative time were approximately 1800 °C and 1000s, respectively.  2.4. Characterization  The fracture morphologies of samples after thermal shock resistance testing and the surface and cross-sectional morphologies of samples after oxidation and ablation testing were observed using the electron microscopy ((SEM, FEI Helios Nanolab 600i, USA) combined with energy dispersive spectroscopy (EDS, INCA Energy 350). The backscattered electron micrograph (BSE, FEI Helios Nanolab 600i, USA) of polished surface of sintered Cf/ZrB2-SiC sample was applied to investigate the distribution between PyC coated carbon fiber and ceramic matrix. The phase compositions of Cf/ZrB2-SiC composite after oxidation testing were detected by the X-ray diffraction (XRD, X′Pert Pro MPD, Holland). The mass ablation rate and linear variation rate of the Cf/ZrB2-SiC composite could calculate by the following expressions,  m  −  St  l  m  0  M  r  =  l  r  =  l  0  −  t  (1)  (2)  Where Mr is the mass ablative rate (mg/mm2·s), m0 and m are the weights of before and after ablation, and the S represents the square of the ablation zone (mm2). The lr presents the line ablative rate (mm/s), l0 and l the heights in the ablation center zone before and after ablation (mm), respectively. And the t presents the ablative time (s).  3. Results and discussion  3.1. Microstructure evolutions  The microstructures of polished surface of Cf/ZrB2-SiC composite were presented in Fig. 1. It was worth noting that a dense and homogenous architecture was achieved (Fig. 1a and b), where the ZrB2-SiC ceramic phases were uniformly distributed between carbon fibers spanning a large macroscopic area. Moreover, the ceramic matrix showed a great sintering characteristic with few porosities, and degradation of carbon fiber was effectively inhibited. Therefore, the dense and uniform architecture with high ceramics sintering characteristic and undamaged carbon fiber could realize the synergistic effect of toughening and anti-oxidation properties. The backscattered-electron (BSE) images of the polished cross-section for Cf/ZrB2-SiC composite with different magnifications were presented in Fig. 2. From the low magnification of polished cross-sectional sample, the fiber bundles in the xy and z direction were uniformly filled with ZrB2-SiC ceramic matrix, showing a uniform distribution between carbon fiber and ceramic matrix spanning a large macroscopic area. Moreover, from the high magnification of polished cross-sectional sample, the ceramic phases were uniformly distributed into the intrafascicular and interfascicular spaces of fiber preform, in which the white phase, dark phase and black phase represented the ZrB2, SiC and carbon fiber, respectively.  3.2. Thermal shock resistance  The thermal shock resistance was an importance index to evaluate the damage tolerance of ceramics, and a high thermal shock resistance represented a superior reliability. The curves between thermal shock temperature difference and residual flexural strength of the Cf/ZrB2-SiC and ZrB2-SiC-Csf composites after water quenching testing were shown in Fig. 3. The residual flexural strengths of the Cf/ZrB2-SiC and ZrB2SiC-Csf composites were keep consistence with the increase of thermal shock temperature difference up to 700 °C, and the residual flexural strength decreased sharply when the thermal shock temperature difference increased up to 800 °C or higher temperature. Moreover, the retained flexural strength of Cf/ZrB2-SiC sample gradually decreased only to 59% of the initial flexural strength at a thermal shock  2  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  Fig. 1. The microstructures of the polished surface for the Cf/ZrB2-SiC composite with different magnification (a) magnification from region C (d) magnification from region D.  low magnification (b) high magnification (c)  temperature difference of 900 °C, exhibiting a high thermal shock temperature difference of 788 °C, remarkably higher than that of ZrB2based ceramics toughened by particles, graphite, graphene or shortchopped carbon fiber (Table 1) [11,25,27-32]. The introduction of continuous carbon fiber could significantly improve the damage tolerance of traditional ZrB2-based ultra-high temperature ceramics. Finite element analysis was used to simulate the temperature and thermal stress distributions in the flexural strength specimen during instantaneous water quenching testing. The un-uniform cooling phenomenon was detected in the surface of the rectangular geometry residual flexural strength bars owing to the higher heat transfer surface area in the corner, resulting in the higher cooling rate in the corner than that in middle surface [33]. The simulated temperature and thermal stress distributions in the flexural strength specimen quenched at the  thermal shock temperature difference of 700 °C and 800 °C were presented in Fig. 4. During the water quenching process at the temperature difference of 700 °C, a maximum temperature difference of 496 °C was detected between the core of specimen (700 °C) and the tensile surface (204 °C). However, the gradient thermal stress between the edges of the test bar (173 MPa) and the tensile surface of the mid-line of the width of the bar (118 MPa) was only approached to 55 MPa, and the maximum tensile stress of 173 MPa at the edges of the bar was lower than the original strength of 204 MPa, indicating that the no flaw or crack could occur in the edge part or tensile surface of testing bars and exhibited no significant strength attenuation, which was consistence with the results of residual flexural strength after water quenching at 725 °C. When the thermal shock temperature difference increased to 800 °C, the thermal stress at the edges of bars was approached to 221 MPa, higher than that  Fig. 2. The backscattered-electron (BSE) images of the polished cross-section of high magnification.  the Cf/ZrB2-SiC composite with different magnification. (a) low magnification (b)  3  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  morphologies in different thermal shock temperature difference were shown in Fig. 5. There was obvious fiber pull-out phenomenon detected when the thermal shock temperature difference was below to 700 °C, which was attributed to the interface debonding mechanism. When the thermal shock temperature difference increased to 800 °C or higher, the fracture surface was flat and carbon fiber pull-out phenomenon was hardly detected owing to the fiber degradation at high thermal shock temperature, resulting in a decrease in residual flexural strength.  3.3. Ablation resistance  To evaluate the oxidation resistance of the Cf/ZrB2-SiC composite, the composite was oxidized in air at 1500 °C for 1 h. The optical images of the Cf/ZrB2-SiC composite before and after static oxidation were presented in Fig. S1. Apparently, there was no cracks detected on the surface of the Cf/ZrB2-SiC composite, and the formed oxide scale (SiO2) was adherent to the oxidized sample. Fig. 6 showed the XRD patterns of the Cf/ZrB2-SiC composite before and after static furnace oxidation. According to XRD analysis, the Cf/ZrB2-SiC composite consisted of ZrB2, SiC, ZrO2 and there was no carbon fiber diffusion peak was detected, which was ascribed to the similar diffusion angles of the carbon fiber and the (001) peak of ZrB2 [11]. Moreover, the ZrO2 diffusion peak might be attributed to the impurity covered on the surface of nanosized ZrB2 powders [14]. After static oxidation in 1500 °C for 1 h, there were some oxidized products like ZrO2 and SiO2 detected on the surface of sample, as a consequence of the oxidation of ZrB2 and SiC phases. No B2O3 peak was detected in the XRD pattern, indicating that a heavy evaporation of B2O3 took place due to the relatively low melting temperature (445 °C) of B2O3 [34]. Moreover, the surface microstructures combined with EDS analysis of oxide Cf/ZrB2-SiC sample were shown in Fig. S2. The surface of Cf/ZrB2-SiC sample after static oxidation was covered by a dense glassy layer, and the dense glassy layer could prevent oxygen from further infiltrating into the material. Moreover, the B2O3 phases produced by the ZrB2 oxidation were apt to break through the dense oxide SiO2 layer due to the capillary force and increasingly vapor pressure, leading to an “island” structure on the partial area of oxide surface. According to the EDS analysis, the island structure consisted of ZrO2 and SiO2 phases, which was consistent with the XRD results.  Fig. 3. The curves between residual flexural strength and thermal shock temperature difference of the Cf/ZrB2-SiC and ZrB2-SiC-Csf composites.  thermal  shock temperature difference (ΔT) with different  Thermal shock temperature difference (ΔT)  Table 1  The comparison of UHTCs.  Materials  ZrB2-SiC ZrB2-nanoSiC ZrB2-SiC-ZrO2 ZrB2-SiC-AlN ZrB2-nanoSiC-graphite ZrB2-SiC-graphite ZrB2-SiC-graphene ZrB2-SiC-Csf Cf/ZrB2-SiC  395 °C 336 °C 461 °C 408 °C 425 °C 583 °C 425 °C 754 °C 788 °C  Reference  [25] [27] [28] [29] [30] [31] [32] [11] This work  of original strength, which could induce the micro crack generation, leading to a decrease in residual flexural strength. In order to reveal the variation tendency in the residual strength of Cf/ZrB2-SiC samples after water quenching testing, the fracture  Fig. 4. The temperature and stress distributions in the specimen calculated by FEA during the water quenching at thermal shock temperature difference of 700 °C (a,b) and 800 °C(c, d).  4  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  Fig. 5. the fracture surface morphologies of the Cf/ZrB2-SiC composite after water quenching testing in different thermal shock temperature.  testing for the Cf/ZrB2-SiC composite. The response time of the sample was approached to 300 s, which was attributed to the high radiation of ZrB2-SiC matrix [38]. The ablation state of the Cf/ZrB2-SiC composite was inset in Fig. 8, the ablation flame color of sample in 800 s was bright than that of sample in 200 s, indicating a higher ablation temperature. The ablation samples kept its shape intact after oxyacetylene ablation for 1000s with maximum temperature exceeding to 1800 °C, as shown in Fig. S3, indicating good ablation resistance of sample exposing high temperature for long time. A layer of whitewash combined with gray oxide was generated on the surface of the Cf/ZrB2-SiC composite after oxyacetylene ablation testing. Moreover, the mass loss rate and line loss rate were 7.32 × 10−4 mg/mm2·s and -1.70 × 10−4 mm/ s, respectively. In order to revel the surface images of Cf/ZrB2-SiC composite after oxyacetylene ablation testing, the inside and outside microstructures of ablation zone were presented in Fig. 9. The outside of ablation zone was covered by the dense oxide scale, which could act as an effective barrier to limit the inward diffusion of oxygen. Moreover, according to the EDS analysis, the dense oxide layer was consisted of the enriched silica and a small amount of zirconia, the enriched silica was attributed to the active oxidation of SiC induced by low oxygen partial pressure, and the zirconia came from the oxidation of ZrB2 [39,40]. In additional, the inside of ablation zone was covered with a glass layer, and there was no penetrative cracks or fiber breakage detected. Moreover, the pores left by the oxidation of carbon fiber or undamaged carbon fiber was covered by the glass layer, which could prevent Cf/ZrB2-SiC composite from being ablated and maintain non-ablative properties. From the results of EDS analysis, the glass layer was composed of large amount of ZrO2 and few SiO2 phases owing to the oxidation of ZrB2 and SiC phase. The ZrO2 phase was uniformly dispersed in the continuous SiO2 phase, as shown in the inset image of Fig. 10b, which could prevent the low melting point SiO2 from being peeled off by the long-time high-velocity and high-pressure oxyacetylene torch [41], greatly improving the ablation resistance of Cf/ZrB2-SiC composite. Fig. 10 presented the microstructures of the cross section of the ablated center sample after oxyacetylene ablation testing. It was worth noting that the oxide layer was composed of multi-layer, and the phases distribution of ablated sample from top to bottom were SiO2-rich oxide  Fig. 6. The XRD patterns of the Cf/ZrB2-SiC composite after static oxidation.  In order to understand the static oxidation behavior of the Cf/ZrB2SiC composite, the cross-section morphologies were presented in Fig. 7. When the static temperature approached to 1500 °C, the molten SiO2 would be forced out from the underneath to the surface due to the capillary force and increasingly vapor pressure and exhibited a super stability in air under 1500 °C, leaving an enriched ZrO2-SiO2 layer (Fig. 7c) underneath the molten SiO2 oxidation layers (Fig. 7b). The enriched ZrO2-SiO2 layer could provide a three-dimensional continuous oxide framework to ensure the configurational stability of the material [35]. In addition, the volume expansion during the transformation of ZrB2 oxidized to ZrO2 could partially fill the residual pores leave by carbon fiber oxidation [24], exhibiting a good oxidation resistance. Finally, two prominent layers were revealed in the cross section of the oxide scale by the EDS scale with a dense SiO2-rich glassy layer (10 μm) and a ZrO2+SiO2 enriched layer (25 μm), respectively, which was also confirmed by previous literatures [36,37]. Fig. 8 presented the temperature rising curve of the oxyacetylene  5  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  Fig. 7. Cross-section micrographs of the Cf/ZrB2-SiC composite after static oxidation at 1500 °C.  the molten SiO2 produced from SiC oxidation would be forced out from the underneath to the surface owing to the effect of capillary force and vapor pressure. The molten SiO2 could be protected to 1850 °C without evaporation for a short oxyacetylene testing time, and the SiO2 layer would be peeled off for a long-time oxyacetylene testing condition. The molten SiO2 out layer would be supplied by the internal SiO2 owing to the capillary force and high vapor pressure, leaving a SiC depleted region layer near ZrO2-SiO2 layer. Therefore, the fabricated Cf/ZrB2-SiC composite exhibited excellent anti-ablation performance under longterm high temperature.  C (s)  1/2O (g)  2  +  CO (g)  →  C (s)  +  O (g)  2  →  CO (g)  2  2ZrB (s)  2  5O (g)  2  +  2ZrO (s)  2  →  B O (l)  2  3  →  B O (g)  2  3   2B O (l)  2  +  3  Fig. 8. The surface temperature rising curve of the Cf/ZrB2-SiC composite after oxgengy ablation.  SiC (s)  +  O (g)  2  SiO (g)  CO (g)  +  →  SiC (s)  2O (g)  2  +  SiO (l)  2  →  CO (g)  2  +  (3)  (4)  (5)  (6)  (7)  (8)  layer, SiO2+ZrO2 oxide layer, SiC depletion layer (ZrO2 layer) and matrix phase according to the EDS analysis. Moreover, the line scanning combined with elements distribution was analyzed to confirm the oxide layer thickness, and the results were shown in Fig. S4 The thickness of rich SiO2 oxide layer, SiO2+ZrO2 oxide layer, ZrO2 layer were 8 μm, 12 μm and 21 μm, respectively. Fig. 11 showed the schematic diagram demonstrating the ablation mechanism of Cf/ZrB2-SiC composite. During the ablation process, the carbon fibers on the surface of sample would react with oxygen to produce CO or CO2 gas when the ablation temperature was up to 400 °C, as shown by chemical Eq. (3) and (4). And the ZrB2 began to oxidize when the temperature rose to 700 °C, forming a liquid-phase B2O3 protective film with antioxidant activity and a porous solid ZrO2 with melting point of 2700 °C, as showed in reaction (5). The B2O3 exhibited a good oxidative protection ability to prevent further ablation of carbon fiber when temperature lower than 1100 °C, and the liquid B2O3 would volatilize rapidly owing to high vapor pressure with the temperature increasing. When temperature increased above 1300 °C,  4. Conclusions  The Cf/ZrB2-SiC composite was successfully fabricated by vibrationassisted slurry impregnation combined with low-temperature hot pressing. The thermal shock resistance and high temperature anti-ablation properties of Cf/ZrB2-SiC composite were conducted by the water quenching method accompanied with stress/temperature filed analysis and oxyacetylene testing, respectively. The thermal shock temperature difference was up to 788 °C, remarkably higher than that of ZrB2-based ceramics toughened by particles, graphite, graphene or short-chopped carbon fiber. Moreover, the thermal stress at the edges of bars was approached to highest, resulting in a decrease in residual flexural strength due to the initiation of micro cracks. After oxyacetylene testing at 1800 °C for 1000s, the mass loss rate and line loss rate were 7.32 × 10−4 mg/mm2·s and -1.70 × 10−4 mm/s, respectively. The low mass loss and line loss rate of Cf/ZrB2-SiC composite were attributed to the enriched silica glass layer embedded with a small amount of zirconia particles. The continuous carbon fiber reinforced ultra-high  6  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  Fig. 9. The SEM and EDS analysis of the surface sample after oxyacetylene ablation testing (a,c) outside of ablation zone (b,d) inside of ablation zone.  temperature ceramics was an important consideration for potential use in thermal protection system applications at high temperature.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to  Fig. 10. The cross section microstructures of oxide layer (a) the overall  layer (b)SiO2 layer (c) SiO2+ZrO2 layer (d)ZrO2 layer.  7  \\x0c', 'D. Zhang, et al.  Corrosion Science 161 (2019) 108181  [17]  [34]  [37]  [21]  [23]  [25]  properties, J. Eur. Ceram. Soc. 39 (2019) 798-805. [15] D. Huang, M.Zhang Q. Huang, L. Wang, L. Xue, X. Tang, K. He, Ablation mechanism of C/C-ZrB2-ZrC-SiC composite fabricated by polymer infiltration and pyrolysis with preform of Cf/ZrB2, Corros. Sci. 98 (2015) 551-559. 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The schematic diagram demonstrating the ablation mechanism of Cf/ ZrB2-SiC composite.  influence the work reported in this paper.  Acknowledgements  This research was financially supported by the National Natural Science Foundation of China (Project Nos. 51872059, 51772061, 51902067 and 11572353), the National Fund for Distinguished Young Scholars (Project No. 51525201), the China Postdoctoral Science Foundation (No.2019M651282), the Major State Basic Search Program (No. 2014CB46505).  Appendix A. Supplementary data  Supplementary material related to this article can be found, in the online version, at doi:https://doi.org/10.1016/j.corsci.2019.108181 .  References  [6]  [4]  [1] N.P. Padture, Advanced structural ceramics in aerospace propulsion, Nat. Mater. 15 (2016) 804-809. [2] D.M. Vanwie, J.R.D.G. Drewry, D.E. King, et al., The hypersonic environment: required operating conditions and design challenges, J. Mater. Sci. 39 (2004) 5915-5924. [3] D. Davidenko, R. Devillers, J. 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},{
  "_id": 138,
  "PDF": "Oxidation behavior and kinetics of ZrB2 containing SiC chopped fibers.pdf",
  "Text": "['Journal of the European Ceramic Society 35 (2015) 4377-4387  Contents lists available at www.sciencedirect.com  Journal  of  the  European  Ceramic  Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Oxidation behavior and kinetics of ZrB2 containing SiC chopped fibers  Laura Silvestroni a,∗ , Elena Landi a , Katarzyna Bejtka b , Angelica Chiodoni b , Diletta Sciti a  a CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy b IIT, Center for Space Human Robotics @Polito, Istituto Italiano di Tecnologia, Corso Trento 21, 10129 Torino, Italy  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  The oxidation of ZrB2 toughened with SiC short fibers has been evaluated in a thermo-gravimetric analyzer from 1200 to 1500   C for 10 h and in a bottom-up loading furnace up to 1650   C for 5-30 min. The microstructure has been  investigated by SEM and TEM. The oxidation kinetics, after a transient stage, is not protective above 1400   C. Fibers follows a  linear behavior,  indicating that the outer oxide  layer  notably oxidize above 1500   C with formation of silica glass and a carbon-rich phase. Mechanisms of SiC fiber oxidation and formation of the various scales are proposed for the various temperature ranges. Comparisons with the oxidation of ZrB2 containing SiC in form of particles, whiskers and previous generation fibers are also discussed.  © 2015 Elsevier Ltd. All rights reserved.  Article history:  Received 15 May 2015 Received in revised form 14 July 2015 Accepted 19 July 2015 Available online 1 September 2015  Keywords:  Oxidation ZrB2 SiC fiber Kinetics Microstructure  1.   Introduction  Zirconium diboride, ZrB2 , is one of the most investigated materials  in the ultra-high temperature ceramics (UHTCs) class due to its  interesting combination of physico-chemical and engineering properties, which places this compound among the most suitable for application  in the hottest parts of next generation aerospace vehicles  [1,2]. The high potential of ZrB2 -ceramics,  in  terms of mechanical and oxidation properties, is approached through addition of secondary phases enabling the achievement of fully dense microstructures and through the introduction of specific dopants, fundamental  for  the high  temperature behavior performances. Among all, the addition of silicon carbide particles has been object of numerous  investigations, owing to their capability to help the densification,  increase  the strength and  toughness and  improve the oxidation resistance  in  the middle-high  temperature regime (1200-1650  C) [1-4]. The kinetics of ZrB2 -SiC composites has been thoroughly studied in several atmospheres, in long and short terms, in various temperature ranges, with different oxidation  facilities, with different amounts of SiC, and accurately modeled [5-11]. ZrB2 preferentially oxidizes in the low temperature range above 600  C, layer only above 1200  C. while SiC starts  to  form a glassy silica  Therefore the first products of the oxidation are a porous ZrO2 scale covered by glassy B2O3 , which, however, has  low vapor pressure and tends to evaporate above 1100  C, as shown by reactions (1)  ∗ Corresponding author. E-mail address: laura.silvestroni@istec.cnr.it (L. Silvestroni).  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.07.024 0955-2219/© 2015 Elsevier Ltd. All rights reserved.  and (2). The  introduction of SiC brings to the formation of a continuous borosilicate glass with higher viscosity and lower oxygen diffusion rate as compared to pure boron oxide [12], reactions (3) and (4), which remains protective up to about 1500  C [6].  ZrB2 +  5/2O2 (g)   →   B2O3 +  ZrO2  B2O3 (l)   →  +  SiC    3/2O2 (g)   →   B2O3 (g)  SiO2 +  CO(g)   SiO2 ·B2O3 (l)   →  SiO2 (l)    B2O3 (l)   +  (1)  (2)  (3)  (4)  However one must take into account that above this temperature, passive-active reactions occur in the scale beneath the outer oxide layer at the SiC-SiO2 interface, due to changes in the partial oxygen pressure, reactions (5) and (6) [13-18].  SiC   SiC   +   SiO2 →  2SiO(g)   +   CO(g)   +   O2 (g)   →   SiO   +   CO(g)   (5)  (6)  It has been  recently  reported  that  the  introduction of SiC  in form of chopped fibers to ZrB2 matrices  increases the toughness by 40%, without jeopardizing the oxidation resistance, especially at high temperature [19,20]. However, the  introduction of SiC fibers instead of SiC particles could alter the mechanisms of oxidation, because of the different spatial distribution of the fibers and their different chemical nature. SiC fibers are nano-composites contain␤-SiC crystals, intergranular carbon and a partially amorphous ing  Si C phase  [21,22]. Therefore  the oxidation of  these  secondary species has to be considered as well.                                              \\x0c', '4378   L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387  The oxidation of sintered SiC fibers belonging to the 3rd generation,  like Hi-Nicalon  S  and  Tyranno  SA3  [22,23], has been extensively studied  in Wright Patterson  labs  in dry and wet air [24-27]. They reported parabolic kinetics in both dry air and laboratory air during the initial oxidation stage at 1000-1200  C, with an activation energy on the order of 180-190 kJ/mol, which  is close to  the value  reported  for  the oxidation of single crystal 6H-SiC [25]. They also observed that crystallization of the amorphous silica to trydimite and crystobalite occurs between 1000 and 1400  C, with a high density of dislocations, cracks and porosity, especially above 1300  C. Evidence of carbon in the amorphous SiO2 was found as well  [27]. The kinetic parameters of  the crystallization of  the amorphous silica were also  found to be strongly sensitive to the impurities and the oxidation environment [27]. The  intense activity carried out so  far on boride-SiC chopped fiber composites [19,20,23] disclosed new basic mechanisms concerning the complex relationships among sintering, microstructure evolution and effects on the mechanical properties, but at the same time,  stirred up new  stimulating  interrogatives  concerning  the thermal behavior. How does the fiber oxidize in various oxidation regimes? Is SiC in fiber shape really more efficient in imparting oxidation resistance than particles? If yes, is there a range in which one type of morphology is preferred over another one? Does the fiber type influence the oxidation behavior of the composite? These are some questions that this study aims to solve. In the present work, the oxidation behavior of ZrB2 containing 3rd generation SiC short fibers was studied by thermo-gravimetric analyses up to 1500  C, to extract the oxidation kinetics, and in a bottom-loading furnace up to 1650  C, to investigate the degradation mechanisms at high temperature. ZrSi2 was selected as sintering additive in order to avoid the  introduction of alien cations and not to modify the chemistry of the already complex system.  2. Experimental procedure  A ZrB2 -based ceramic was prepared using the  following commercial powders: hexagonal ZrB2 , Grade B (H.C. Starck, Germany), specific  surface area 1.0 m2 /g,  impurity max  content  (wt%): C: 0.25, O: 2, N: 0.25, Fe: 0.1, Hf: 0.2, particle size range 0.1-8  \\u242em; orthorhombic ZrSi2 -F (Japan New Metals Co., LTD, Osaka,  Japan) particle size 2.8  \\u242em,  impurities  (wt%): C: 0.11, Fe: 0.09, O: 1.00; Tyranno SA3 SiC fibers (UBE Europe GmbH, Dusseldorf, Germany) with composition (wt%) Si: 67.8, C: 31.3, O: 0.3, Al < 2, diameter: 7.5  \\u242em, chopped length: 100-200  \\u242em. A composite containing 10 vol% of fiber and 10 vol% of zirconium silicide was prepared by conventional wet milling route and sintered by hot pressing to full density, as explained  in details  in [23]. To study  the oxidation kinetics,  thermo-gravimetric analyses were  carried out  at  atmospheric pressure on  specimens with dimensions 8 mm   8 mm   2 mm  in dry  synthetic  air  (80 vol% N2 + 20 vol% O2 , with 30 ml/min gas flow), in the temperature range of 25-1500  C, with  isothermal exposure  times of 10 h, heating rates of 30  C/min and free cooling (model STA449, NETZSCH, Geraetebau GmbH, Selb, Germany). The mass variation was recorded continuously with 1  \\u242eg of sensitivity. Parallel oxidation  tests were carried out  in a bottom-loading furnace (Nannetti FC18, Faenza,  Italy and Carbolite BLF1800  furin static air at 1500 and 1650  C  nace, Hope Valley, UK)  for 5, 15 and 30 min on  rectangular 13 mm   2.5 mm   2 mm bars. Specimens were located on porous zirconia supports in the furnace when the maximum temperature was achieved and then removed and air-quenched after the exposure time. The mass and dimensions of the bars were measured before and after the oxidation. The microstructures of  the as-sintered material and oxidized specimens were  analyzed  using  field  emission  gun  scanning  ×  ×  ×  ×  electron microscopy  (FE-SEM, Carl Zeiss Sigma, Germany)  and energy dispersive  spectroscopy  (EDS,  INCA Energy 300, Oxford Instruments, UK) on  surface and  fractured  cross-section of  the specimens, to reveal the modifications induced by oxidation. Transmission electron microscopy (TEM) analyses were performed on the as-sintered material and on selected oxidized samples. A TEM specimen  from bulk ceramic was prepared by cutting 3 mm disk from the sintered pellet. This was mechanically ground down to about 20  \\u242em and then further ion-beam thinned until small perforations were observed by optical microscope. For the analysis of a selected oxidized sample, a lamella was extracted from an area of interest using a focused ion beam (FIB, Carl Zeiss Dual Beam Auriga, Germany). The final  thinning was performed with  low currents and  then  the  lamella was cleaned and polished  in conditions of low current and low ion beam energy conditions. TEM images and EDS analyses were acquired using a transmission electron microscope  (TEM, FEI Tecnai F20 ST, The Netherlands) equipped with a Schottky emitter and operating at a nominal voltage of 200 keV, combined with a EDAX EDS X-ray liquid nitrogen cooled spectrometer PV9761/II with a super ultra-thin window.  3. Results and discussion  3.1. As sintered microstructure  The ceramic was fully densified by hot pressing at 1600  C. ZrB2 mean grain size was around 2.5  \\u242em and  the fibers were homogeneously dispersed without agglomerates, as put  in evidence by the SEM  image  in Fig. 1a. A magnification of  the fiber section  is shown  in Fig. 1b:  it has a rounded shape with smooth  interface. Carbon pockets are homogeneously distributed all over the fiber up to about 200 nm  from the fiber edge, where dense SiC grains comprise a  thin  rim. The mean grain  size of  the crystallites  in the core  is around 75 nm, while  the grain size of SiC  in  the rim increased up to 245 nm during sintering [23]. TEM  images of the fiber  in Fig. 1c and d evidence the core-rim regions and show the nature of the  intergranular phase, containing turbostratic carbon and amorphous Si  C, Fig. 1e. At  the fiber-matrix  interface, secondary phases  like ZrO2 , ZrSi2 or Zr  Si  C  O were often  found, as shown  in Fig. 1b and  in the EDS spectra of Fig. 1, however the interfaces between fiber-ZrB2 , fiber-ZrSi2 and adjacent ZrB2 were generally clean, Fig. 1f. This material showed a good combination of mechanical properties with 4-point bending strength of 465 MPa and improved fracture toughness of 5.23 MPa m, as compared to the unreinforced ceramic, 4.25 MPa m [23,29].  √  √  3.2. Oxidation in TG  3.2.1.   TG curves  The typical heating curve up to 1500  C, preceding the isothermal stage, is depicted in Fig. 2a, while TG curves measured during to 1500  C  isothermal holds  from 1200  for 10 h  are  shown  in Fig. 2b. During heating, Fig. 2a, after a  rapid mass  increase  in the first 10 min, a slower  increase occurs with temperature up to 800  C,  then a slightly higher  increase  is observed between 800 and 1200  C,  followed by a new  increase at higher rate  for  temperatures above 1200  C. According to similar profiles observed for other ZrB2 -based materials [30], air oxygen is chemisorbed on the surface of the sample and then adsorption-desorption equilibrium is reached at around 200  C, Fig. 2a. The formation of solid products, described by reaction (1), slowly progresses up to 800  C. Oxidation of the main secondary component (SiC), described by reaction (3), begins at around 1000  C and then steadily proceeds up to 1200  C, with continuous loss of volatile species, mainly B2O3 , expressed by reaction (2). In the following stage of interaction between sample  \\x0c', 'L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387   4379  Fig. 1. Microstructure of the as sintered composite showing (a) the fiber distribution, (b) the fiber section, (c) the magnified view of the fiber morphology evidencing, (d) the coarsened rim, (e) the nature of the intergranular C-rich pockets and (e) clean fiber/matrix interface and in the bottom EDS of the secondary phases in (b).  Fig. 2. TG curves recorded during (a) the heating stage, (b) the isothermal hold for 10 h at 1200-1500   C and (c) sketch of the curve shape during isothermal hold with the main phenomena taking place.  to 1400  C), a stable amorphous borosilicate and air  (from 1200  glassy film forms on the surface, described by reaction (4). The last weight gain increase, from 1400 to 1500  C, is due to oxygen penetration through the bulk and to formation of new solid oxide phases.  During the isothermal hold, Fig. 2b, all the curves display a hump at the beginning of the dwell: the weight gain tends to  increase, and then, after a certain time,  it decreases  indicating the  leakage of gaseous species. After  that,  it  increases again almost  linearly,  \\x0c', '4380   L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387  Fig. 3. SEM images of the surface after TG analysis at (a-c) 1400   C, (d-g) 1500   C and 10 h dwell. (a and d) Surface appearance, (b and e) morphology of the ZrO2 precipitates in the glass, (c and f) fiber aspects, (g) ZrO2 grain morphologies. (h-l) EDS spectra of the phases as indicated.  representing the continuous formation of solid products. In detail, at 1200  C the weight loss is recorded between the 2nd and the 4th hour, and remains stable until the end of the test, suggesting that at this temperature there is a balance between weight loss and weight gain phenomena. At 1300  C the first weight  increase  is recorded until the 1st hour, then the weight gain decreases until the 5th hour, and then it starts to slowly increase again. At 1400  C the duration of the first weight  increase  is only 30 min, than  increases. At 1500  C the weight it decreases until the 4th hour and then  it  linearly  increase hump is further reduced to 20 min, then a linear behavior dominates the oxidation process. A sketch of the behavior displayed by the weight gain is shown in Fig. 2c and will be further discussed in the following sections.  Table 1 Thickness of all the specific layers after TG analyses in the 1200-1500   C range for 10 h dwell.  T,   C   Layer I SiO2  \\u242em   6  11  20  25   1200  1300  1400  1500   Layer II ZrO2 -SiCOfiber pullout  \\u242em   38  63  60  40   Layer III ZrO2 -ZrB2 -SiC  Total thickness  \\u242em   50  50  50  65   \\u242em  94 124 130 130  3.2.2. Microstructure evolution after TG  The specimens after TG tests at 1200 and 1300  C (not shown) were all covered by a continuous silica-based glass with nano-sized ZrO2 precipitates, which  tended  to agglomerate and  form round islands with  increasing  temperature, especially at  the specimen edges. Starting from 1400  C, in correspondence of bigger agglomerates, cracks developed  in the glass, as put  in evidence by SEM analyses  in Fig. 3a, owing  to  the  large mismatch between  the  10−7 / C thermal expansion coefficients of  silica and ZrO2 , 5.6   10−6 / C [32], respectively. In addition, some bubbling [31] and 10  occurred, leaving the subsurface layer exposed. Oxidized fibers are recognizable among ZrO2 grains in Fig. 3c. EDS analysis revealed Si, O and C in the fiber core and C-rich Si  O  C in the outer part, Fig. 3h,  ×  ×  with Si:O:C atomic percentage of 14.9:60.2:24.9 and 4:8.5:87.5, in  the core and outer part, respectively. After heat  treatment at 1500  C,  large bubbles were  found on  the  surface, as  shown  in Fig. 3d, and ZrO2 precipitates assumed a dendritic shape preferred over a rounded one, probably owing to Al contaminations which favored anisotropic grain development  [33], Fig. 3e and  i. Under the exploded bubbles,  the fiber sites were filled with silica and tiny Al  Si  Zr  O whiskers, Fig. 3f and  l, while ZrO2 grains had a polyhedral shape, Fig. 3g. The cross section views revealed the same scales structure with variable thicknesses, depending on the oxidation temperature;  in Table 1 the thickness of all the specific  layers after TG analyses  is reported.  \\x0c', 'L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387   4381  Fig. 4. SEM images of the fractured cross section after TG analysis at 1300   C and 10 h dwell. (a) Overview, (b) layer II, (c) magnification of the layer II/layer III interface and (d) layer III. In the bottom EDS of the phases as indicated.  As illustrative example, the scale morphology after TG at 1300  C is shown in Fig. 4a. An outermost compact silica scale (I) tops a ZrO2 layer where the fibers are partially oxidized and pulled out from the matrix owing to some porosity in the oxide and to the formation of a thin C coating around the fibers (II), Fig. 4b. Then a partially oxidized scale follows (III): it is distinguishable from the bulk because of the transgranular fracture, compared to the intergranular observed in the bulk, Fig. 4c. In this region, small ZrO2 grains were found along the ZrB2 grain boundaries, and the fibers were almost unoxidized, Fig. 4d. It is interesting to note in Table 1 that the thickness of SiO2 increases with increasing temperature, while that of the subsurface layers is non-linear with the temperature; in particular, the thickness of layer II tends to be stable at 1500  C. This is probably due to lower boron oxide content in the outermost glassy layer, which is continuously impoverished in B2O3 owing to its faster vaporization at 1500  C [10]. In these conditions, due to the higher viscosity of silica, the oxygen diffusion coefficient  is  lower and oxygen penetrates a thinner subsurface scale, resulting  in a total thickness of modified material at 1500  C even thinner than to that of samples oxidized at 1300 and 1400  C.  3.2.3. Oxidation kinetics in the 1200-1500   C range  The TG curves  follow similar trends to those outlined by previous works on ZrB2 -SiC ceramics with SiC  in  form of particles [19,34-36] and a sketch of  the occurring events  is presented  in Fig. 2c. A hump  is visible  in all  the curves and  is similar  to  the mechanisms occurring  for oxidation of ZrB2 -SiC particles,  so  it can be deduced  that  in  regime 1, ZrB2 oxidizes  faster  than SiC and forms solid zirconia and liquid boron oxide, resulting in a net weight increase. The drop of weight gain is then due to boron oxide  Fig. 5. Arrhenius plot of the oxidation rate constants, k, against 1/T assuming linear kinetics.  evaporation, regime 2, reaction (2). The hump amplitude indicates the duration time for the evaporation of boron oxide, before entering in a nearly stable state. In regime 3, the weight  loss is slowed down by the oxidation of SiC fiber, which  leads to the  formation of silica and then of a borosilicate glass. As  long as more boron  is dissolved in the silica-based glass, it becomes less viscous and thus enables the penetration of oxygen through the scale  [37], which loses its protective action. Therefore, after a certain period, a following weight increase takes place, regime 4, owing to the formation of higher amount of solid products  (ZrO2 , SiO2 , C), which overpasses the kinetics of evaporation of gaseous species (B2O3 , SiO, CO). Notably, the kinetics of regime 4 never suggests a parabolic or  \\x0c', '4382   L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387  Fig. 6. Oxidation mechanism in ZrB2 -SiCf composite (see also Fig. 2c for weight gain regimes).  paralinear  behavior, which would  correspond  to  a  protective behavior,  as  claimed  for  the  addition of  SiC particles  to ZrB2 [19,34,38]. By considering the oxidation curves  in the  isothermal stage, Fig. 2b, i.e. up to 360 min at 1200 and 1300  C, up to 230 min at 1400  C and up to 20 min at 1500  C, the kinetics follows a discontinuous trend, which evolves to a  linear one, representing an interface controlled oxidation caused by the breaking of the continuous oxide scale  formed during  the  initial stage of oxidation, in agreement with the bubbles and cracks observed  in the external silica-based layer [39,40]. After 6 h for all temperatures, we can assume a linear kinetics behavior which follows the linear law:  \\x01W  S  =   kT  (7)  where  \\x01W/S is the mass change of the sample per unit area, T is the oxidation temperature and k is the oxidation rate constant. This last parameter was calculated for all TG temperatures and the resulting goodness at R2 was higher than 0.985. In Fig. 5, the Arrhenius plot of k against temperature is reported. The apparent activation 253 kJ/mol, that energy, Ea, derived from an Arrhenius equation is  is higher than values reported for ZrB2 -SiC particles [38] or solely SiC particles and Hi-Nicalon S fibers [27]. According to the microstructural features of the specimens after TG and the weight-gain curves, the evolution of the sub-surfacial scales and the role of SiC fibers on the oxidation process is proposed and simplified  in the sketch of Fig. 6. At the beginning of the oxidation, i.e. regime 1 and 2, mainly ZrB2 is involved, as described by reactions (1) and (2). From regime 3, fibers start to oxidize following a passivating mechanism and to combine with ZrB2 oxidation products, expressed by reactions (3) and (4). The observation of carbon together with silica around the fiber indicates that oxidation of SiC fibers in the subsurface layer is different than the one described by reaction (3), and could occur as follows:  SiC   +   O2 (g)   →   SiO2 +   C   (8)  in agreement to what found  in previous studies on the oxidation of Hi-Nicalon fibers  in ZrB2 [20,21]. Thermodynamics  in standard \\x01G  equal to  −541 kJ conditions give reaction (8)  favorable with  at 1500  C. Entering  in regime 4,  in the subsurface  layer,  further silica is released by the SiC oxidation, together with CO developed by oxidation of C. The tested conditions can be considered mild, as the maximum temperature is 1500  C and the amount of oxygen in the TG apparatus is limited, so the architecture of the final specimens does not include the formation of a SiC-depleted ZrB2 scale.  Table 2 Comparison of weight  facilities.  Material   ZrB2 -10 SiCf-10 ZrSi2  ZrB2 -20 SiC-5 Si3 N4 [34]  ZrB2 -20 SiC [10]   ZrB2 -20 SiC [11]  ZrB2 -10.7 SiC-8 B4 C [36]  ZrB2 -21.9 SiC-8 B4 C [36]   gain   among ZrB2 -SiC   composites   tested   in   various TG  TG conditions T (  C), dwell (h)  Apparent weight gain, mg/cm2  1200,10  1300,10  1400,10  1500,10  1120,10  1200,2  1400,2  1400,10  1500,10  1500,1.5  1500,1.5   0.9 0.2 5.8 12.6 0.8 5 5 13.8 10.5 18 8  A direct comparison of oxidized thickness or weight gain among data reported  in the  literature on ZrB2 -SiC composites cannot be very precise due to slightly different experimental conditions, but some values on ZrB2 composites containing SiC particles and oxidized  in TG  facilities are reported  in Table 2. The weight gain of ZrB2 -composites containing SiC fiber or particles,  is comparable till 1200  C  [34] and this value remains almost unchanged up to 1300  C for the fiber-containing material. Data collected at higher temperatures are not always  in agreement. For example, Opeka et al. found that the weight gain increased already at 1200  C, but did not change up to 1400  C for 2 h dwell [10]. It can be also noticed that oxidation at 1400  C for 2 h of ZrB2 -20 SiC particles gave similar weight gain as ZrB2 -10 SiC fibers at the same temperature, but after 10 h. On the other side, in a more recent work by Zapata-Solvas et al. [11], the weight gain at 1400  C was twofold than what found in this study and remained basically constant up to 1500  C, that is in the same order of this ZrB2 -SiC fiber composite. Lastly, in the work of Peng et al. [36], the oxidation at 1500  C for 90 min resulted in higher weight gain than the composite with fiber after dwell for 10 h. From  these  data  it  can  be  deduced  that,  at  least  in  the 1400-1500  C  temperature range,  the oxidation of ZrB2 containing 10 vol% crystalline SiC fibers  is accompanied by  lower weight gain than ZrB2 containing 10-20 vol% SiC in form of particles, probably due to a slower oxidation kinetics of the fibers and to their discontinuous distribution throughout the matrix. The better oxidation resistance as compared to the materials listed in Table 2 is achieved even without the introduction of third phases, like Si3N4 [34] or even B4C [36], which could further slow down the oxidation processes.    \\x0c', 'L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387   4383  Fig. 7. Appearance of the as sintered (Ref.) and oxidized specimens after oxidation in the bottom-loading furnace for 5, 15 and 30 min at (a) 1500   C and (b) 1650   C. (c) Plot of weight gain per surface unit as a function of the oxidation time and temperature.  Fig. 8. Microstructure of the ceramic upon oxidation at 1500   C for 15 min showing (a) the external surface, (b) the cross section, (c) a magnification of a fiber in layer II and (d) EDX spectra of areas indicated in (b).  3.3. Oxidation in bottom-loading furnace  3.3.1. Weight gain and microstructure evolution  The appearance of the bars oxidized for 5, 15, 30 min at 1500 and 1650  C is shown in Fig. 7a and b. It can be noticed that after 15 min at 1500  C, the color gradually changed from gray to dark gray and after 15 min at 1650  C the surface was brighter and macroscopic bubbles appeared. The weight gain per  surface area,  increased quickly in the first 5 min and then increased with time with a slower rate, owing to the formation of solid oxides, as shown in Fig. 7c. At 1650  C the weight  increase  is more sluggish than what observed at 1500  C, probably because of the faster escape of volatile species, like SiO, CO. To remark that  in no case the specimens attached to the support.  Table 3 Thickness of all the specific layers after bottom-loading oxidation tests at 1500 and 1650   C for 5-30 min.  T,   C   Dwell, min  Layer I SiO2 ,  \\u242em  Layer II ZrO2 -SiCOfiber pullout, \\u242em  Layer III ZrO2 -ZrB2 -SiC, \\u242em  Total thickness, \\u242em  1500  1650  5  15  30   5  15  30   a SiC depleted.  4  7  9   15  18  24   26  29  42   45  40  66   14  25  34   27a 36a 57a  44 61 85  87 94 147  \\x0c', '4384   L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387  Fig. 9. Microstructure of the ceramic upon oxidation at 1650   C for 15 min showing (a) the external surface, (b) the cross section and (c and d) magnified view of fibers at the oxide/boride interface (II/III).  The external surface of  the oxidized specimen was homogeneously covered by silica glass after 5 min at 1500  C. After a dwell at this temperature for 15 min, ZrO2 precipitates started to be visin Fig. 8a. After oxidation at 1650  C, the ible, as put  in evidence  surface was completely covered by glassy bubbles and the whole surface became rougher after a dwell of 15 min, as shown in Fig. 9a. The cross sections displayed a  layered architecture with some common features. The thickness of each layer varied depending on \\u242em at 1500  C to 75-150  \\u242em at 1650  C, the dwell time, from 30-50  as shown in Table 3. The growth of the oxide layer exhibited a linear behavior,  in agreement with TG analyses. For simplicity,  the microstructural features of only some selected specimens, oxidized for 15 min, will be illustrated. The cross section of the specimen oxidized at 1500  C for 15 min shown in Fig. 8b shows an outermost continuous silica-based layer (I) adherently attached to a ZrO2 layer partially filled with silica glass (II). Columnar grains are at the early stage of development and extensive fiber pullout takes place in this scale owing to the weakening of the fiber/matrix interface and to the loose microstructure with weak grain boundaries. SiC fibers are evidently corroded at this temperature: the inner part is composed of SiC, but moving outwards the fiber, the main constituents are silica and carbon, Fig. 8c, as confirmed by the EDS spectra in Fig. 8d. EDS on bright particles of the fibers revealed zirconium, boron and Si, C, O. A transition region with SiC and ZrB2 surrounded by zirconia grains  followed (III), and underneath this scale, unoxidized bulk was found. A SiCdepleted  layer was not observed at this temperature,  for neither dwell periods of 15 nor 30 min. Conversely, oxidation at 1650  C was effective  in  initiating the regime of SiC active oxidation  in  the  subsurface, as  in all  the  three samples oxidized at this temperature, SiC-depleted ZrB2 was observed, Table 3. The morphology of the cross section after 15 min at 1650  C is shown in Fig. 9b. The outer 58  \\u242em was composed by a compact SiO2 scale (I), topping ZrO2 grains with elongated shape and veined by silica-based glass (II).  In scale (II), no holes related to fiber consumption were found, owing to the volume expansion occurred by oxidation of ZrB2 to ZrO2 and by ZrO2 grain coarsening. On the contrary, at the ZrO2 -ZrB2 interface (III), cavities and oxidized fibers were observed. Fiber detachment form the matrix was often observed at  this  temperature, as shown  in Fig. 9c.  In addition, a more pronounced progressive fiber oxidation to silica and carbon was  found, Fig. 9d. From  region  (III) a  thin  lamella was extracted using the FIB, as shown  in Fig. 10a. TEM analyses confirmed  the  formation of SiO2 both at  the fiber/matrix  interface, Fig. 10b and c, and at SiC grain boundaries  inside the fiber, Fig. 10d-f.  3.3.2. Degradation mechanisms of SiC fiber in ZrB2 matrix up to 1650   C  The oxidation action in such type of furnace is more aggressive than  in  the previously presented TG experiments, due  to higher oxygen partial pressure, that accelerates all the degradation mechanisms. However, up to 1500  C, the morphology of the various oxide scales follows the scheme described by regimes 1-4 in Fig. 6. At 1650  C, under  the outer silica scale, ZrO2 and only silica residuals are  found  in place of  the original SiC fibers. Concurrent phenomena take place: SiC nanocrystals composing the fiber oxidize and  form SiO2 and gaseous CO, similarly to SiC particles, according to reaction (3), SixCy amorphous phase  inside the fiber  \\x0c', 'L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387   4385  (a) Low magnification FE-SEM image of the thin lamella extracted by FIB from layer (III) of the specimen oxidized at 1650   C for 15 min. (b, c, e) TEM images of the Fig. 10.  lamella showing formation of SiO2 at the fiber-matrix interface or (d and e) among SiC grains in the fiber. (f) EDS spectra of SiO2 overlapped to that of a SiC grain.  Fig. 11. Structure evolution of the SiC fiber in the composite in the various layers during fast oxidation. (a) In the outermost columnar ZrO2 scale, (b) moving inwards in the ZrO2 scale, (c) in the SiC-depleted ZrB2 scale only from 1650   C and (d) in the first ZrB2 bulk.  decomposes as described by reaction (9) [25,26] and subsequently oxidize to gaseous products (SiO, CO):  Six Cy →  xSi   +  yC   (9)  In addition, in layer II of Fig. 9b, abrupt gas escape drive ZrO2 crystal growth along parallel channels, resulting  in columnar grains, irrespective of the original SiC shape or distribution. As SiO vapour upstreams  the ZrO2 scale, oxygen partial pressure  increases and condensation of SiO2 is favorable, according to (10):  SiO   +   1/2O2(g) →   SiO2  (10)  At the interface between layer II and layer III (Fig. 9b-d), due to very low oxygen partial pressure, conditions for active SiC oxidation  are satisfied and active oxidation of the fibers starts, similarly to what happens during oxidation of ZrB2 -SiC composites [5,6]. At  the border between  the  layer  III and  the unaffected bulk (Figs. 9b and 10c), an early stage of fiber oxidation can be observed, with the formation of silica around the fiber and between grains. A sketch of this final stage of the whole oxidation mechanism for the ZrB2 -SiCf composite  is displayed  in regime 5 of Fig. 6. From reactions (3), (5) and (6)  it  is clear that the oxidation of SiC fiber is accompanied by abundant generation of gaseous CO and SiO species.  If  the  leakage of  these gases  is hindered by a compact ZrO2 -SiO2 oxide, fiber sites are saturated in CO which depresses the rate of SiC oxidation reaction [41] allowing the survival of carbon. Fig. 11 summarizes the structure evolution of SiC fiber  in the composite  in the various  layers during  fast oxidation: under the      \\x0c', '4386   L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387  Table 4 Comparison of weight gain of ZrB2 -SiC composites tested in our lab in a bottom-load furnace under the same experimental conditions. T: Tyranno SA3 fiber, H: Hi-Nicalon fiber, p: particle, w: whisker.  Label   ZZ10-T ZS10-H [42]  ZS10-w [43]  ZS20-p [43]   ZZ10-T ZS10-H [42]   Oxidation temperature and dwell,   C, min  Sintering additivevol%   SiC content vol%  SiC form   1500, 30  1650, 30  10 ZrSi2 5 Si3 N4 5 Si3 N4 5 Si3 N4  10 ZrSi2 5 Si3 N4  10  10  10  20   10  10   Tyranno SA3  Hi-Nicalon  Whiskers  Particles   Tyranno SA3 Hi-Nicalon   SiC size \\u242em  Ø: 7.5, l:180  Ø: 15, l:180  Ø: 0.6, l:30  0.6   Ø: 7.5, l:180  Ø: 15, l:180   \\x01wt/S %/cm2  0.014 0.022 0.010 0.024  0.019 0.032  outermost silica scale, among columnar ZrO2 scale, the fiber is completely converted to SiO2 and SiO (Fig. 11a); moving inwards in the ZrO2 scale the fiber is surrounded by C and SiO2 at 1500  C (Fig. 11b). At 1650  C, when active oxidation of SiC starts, holes left by the fiber are found in ZrB2 (Fig. 11c), but in the first ZrB2 zone under ZrO2 , an early oxidation stage of the fiber to SiO2 and C is observed, as at this temperature SiC oxidizes faster than ZrB2 (Fig. 11d).  3.3.3.   Comparison with other ZrB2 -SiC composites  A  direct  comparison  of  oxidized  thickness  or weight  gain among the data reported in the literature on ZrB2 -SiC composites is meaningless owing  to  the different atmosphere,  furnace  type and set-up, specimen shape and size. Therefore we will  limit our considerations  to analyses performed  in our  laboratories  in  the same bottom-loading furnace [42-44]. Table 4 reports the specific weight gain of various ZrB2 composites  containing SiC  in  form of whisker, particle or previous generation Hi-Nicalon SiC fiber [22]. Unfortunately at present no data are available on the analogous material containing 10 vol% SiC in  the  form of particles, but previous studies  [43]  revealed  that whiskers had  the same weight gain as particles  in  the  temperature regime considered. We also have to notice that the sintering additives  in Table 4 are different, and Si3N4 has better oxidation resistance than ZrSi2 because of the nitrogen dissolved in the silicabased  scale,  resulting  in a  tighter glass network  [44]. This was also confirmed by previous studies, where Hi-Nicalon SiC chopped fibers-ZrB2 composites with addition of different sintering aids were oxidized at 1650  C for 30 min  in same furnace [45].  In that work, ZrSi2 was confirmed to negatively affect the overall oxidation resistance of the composite, especially  in comparison with Si3N4 , mainly due to the low melting point of ZrSi2 or to secondary Zr-Si phases formed during sintering. Albeit slightly different compositions, we can say that the specific weight gain after oxidation at 1500  C  for 30 min has  little variations, in the 0.010-0.024%/cm2 range, with whiskers displaying the  lowest weight gain. Better dispersion of SiC,  i.e. small SiC size,  results  in quicker  formation of a protective silica glass on the  surface. Besides, ZrB2 containing 10 vol% Tyranno fiber has even better oxidation resistance at 1500  C than ZrB2 sintered with a more  refractory  sintering additive and containing 20 vol% SiC particles  [43], when normally higher SiC content has better performances up to 1650  C [40,46]. The lower weight gain of 10 vol% Tyranno SA3 in comparison to 20 vol% SiC particles can be related to small amount of Al present  in the fiber, that  increases the viscosity of the silica-based glass [24,28]. Therefore, we can suppose that the introduction of the proper sintering additive would further improve the oxidation behavior of ZrB2 -Tyranno SA3 composites. With  regard  to  the fiber  type,  Tyranno  SA3 fiber  is more oxidation-resistant  than Hi-Nicalon,  thanks  to  lower amount of oxygen and amorphous SixCyOz in the initial fiber microstructure. This oxygen-containing amorphous phase is responsible for the formation of pores, arising from CO escape according to the reaction  of decomposition (11) and oxidation (12) of the unstable phase, as previously reported by Takeda et al. in a study about the oxidation of Hi-Nicalon S fiber after exposure for 1-100 h at 1000-1500  C in air [47].  Six CyOz →  xSiO2 +   (z  −   2x)CO(g)   +  −  z  (y  +   2x)C   Six CyOz +   (x  +   (y  −  z))O2 (g)   →  xSiO2 +  yCO(g)   (11)  (12)  Better oxidation behavior of Tyranno SA3 over Hi-Nicalon fiber  is observed in the present study as well up to 1650  C, Table 4 [42]. Therefore, even under fast oxidation  in the temperature range of 1500-1650  C, the present ZrB2 -based composite with 10 vol% of crystalline SiC fibers has comparable weight gain to ZrB2 containing 10 vol% of whiskers, despite the different spatial distribution and size, and  lower weight gain  than ZrB2 containing 10 vol% of previous generation SiC fiber, thanks to a lower oxygen content in the crystalline fiber.  4. Conclusions  The oxidation behavior of a  tough ZrB2 -based ceramic containing  10 vol%  short  SiC  Tyranno  SA3  fiber was  studied  by thermo-gravimetric analysis up to 1500  C for 10 h and in a bottomloading furnace up to 1650  C for 5-30 min. After a transition period  in the very first oxidation stage,  linear kinetics prevailed, unlike the parabolic one observed for ZrB2 or ZrB2 -SiC composites. The apparent activation energy  for  this composite was calculated to be around 253 kJ/mol, that  is higher compared to previously reported activation energy for ZrB2 composites with SiC in form of particles. The weight gain after fast-oxidation in bottom-loading furnace matched well with that of other ZrB2 compounds containing SiC in form of particles, whiskers or previous generation fibers. 10 vol% Tyranno SA3 fiber imparted better oxidation resistance to ZrB2 than 20 vol% SiC particles or 10 vol% Hi-Nicalon fiber. Although the discontinuous SiC distribution, columnar ZrO2 and SiC-depleted ZrB2 from 1650  C, where  formed, but only starting  fast evolution of gaseous CO and SiO occurred, due to the oxidation of the fiber into C and SiO2 . Introduction of 3rd generation SiC fiber to a ZrB2 matrix seems a promising strategy to obtain composites with improved failure tolerance and good oxidation resistance. The SiC fibers provide higher toughness and behave at  least as good as particles  in ZrB2 matrix during oxidation up to 1650  C.  Acknowledgements  The authors wish to thank D. Dalle Fabbriche for oxidation tests in bottom-loading furnace. TEM analyses were carried out through grant N. FA8655-12-1-3004, provided by US Air Force Research Laboratory with Dr. Ali Sayir as contract monitor.                        \\x0c', 'L. Silvestroni et al. / Journal of the European Ceramic Society 35 (2015) 4377-4387   4387  References  zirconium of  oxygen,  [1] S.Q. Guo, Densification of ZrB2 -based composites and  their mechanical and physical properties: a review, J. Eur. Ceram. Soc. 29 (2009) 995-1011. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [3] F. Monteverde, A. 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Soc. 91 (2008) 1475-1480. [20] D. Sciti, L. Silvestroni, Processing, sintering and oxidation behaviour of SiC fibers reinforced ZrB2 composites, J. Eur. Ceram. Soc. 32 (2012) 1933-1940. [21] D. Sciti, R. Savino, L. Silvestroni, Aerothermal behavior of a SiC fiber-reinforced ZrB2 sharp component  in supersonic  regime,  J. Eur. Ceram. Soc. 32  (2012) 1837-1845. [22] T. Ishikawa, Advances in inorganic fibers, Adv. Polym. Sci. 178 (2005) 109-144. [23] L. Silvestroni, D. dalle Fabbriche, D. Sciti, Tyranno SA3 fiber-ZrB2 composites. Part I: microstructure and densification, Mater. Des. 65 (2015) 1253-1263. [24] R.S. Hay, G.E. Fair, R. Bouffioux, E. Urban, J. Morrow, A. Hart, M. Wilson, Relationships between fiber strength, passive oxidation and scale crystallization kinetics of Hi-NicalonTM -S SiC fibers,  in: D. Singh,  J. Salem, S. Widjaja (Eds.),  [19]   [35]   Mechanical Properties and Performance of Engineering Ceramics and Composites VI: Ceramic Engineering and Science Proceedings, vol. 32,  John Wiley & Sons, Inc., Hoboken NJ, USA, 2011, pp. 39-51. [25] P. Mogilevsky, E. Boakye, R. Hay, J. Welter, R. Kerans, Monazite coatings on SiC fibers. II. Oxidation protection, J. Am. Ceram. Soc. 89 (2006) 3481-3490. [26] R.S. Hay, G.E. Fair, R. Bouffioux, E. Urban,  J. Morrow, A. Hart, M. Wilson, HiNicalonTM -S SiC fiber oxidation and scale crystallization kinetics, J. Am. Ceram. Soc. 94 (2011) 3983-3991. [27] R.S. Hay, G.E. Fair, A. Hart, S. Potticary, R. Bouffioux, Kinetics of passive oxidation of Hi-Nicalon-S SiC fibers in wet air: relationships between SiO2 scale thickness, crystallization, and fiber strength,  in: D. Singh, J. Salem, M. Halbig, S. Mathur (Eds.), Mechanical Properties and Performance of Engineering Ceramics and Composites VII, John Wiley & Sons, Inc., Hoboken NJ, USA, 2013, pp. 253-260. [28] Y. Wang, L. Luo,  J. Sun, L. An, ZrB2 -SiC(Al) ceramics with high resistance  to oxidation at 1500   C, Corros. Sci. 74 (2013) 154-158. [29] L. Silvestroni, D. Sciti, C. Melandri, S. Guicciardi, Tyranno SA3 fiber-ZrB2 composites. Part II: mechanical properties, Mater. Des. 65 (2015) 1264-1273. [30] V.O. Lavrenko, A.D. Panasyuk, O.M. Grigorev, O.V. Koroteev, V.A. Kotenko, High temperature oxidation of ZrB2 -MoSi2 ceramics  in air, Powder Metall. Met. Ceram. 51 (2012) 102-107. [31] B. El-Kareh, Fundamentals of Semiconductor Processing Technologies, Kluwer Academic Publishers, Dordrecht, The Netherlands, 1995. [32] R. Stevens, Engineering properties of zirconia, in: S.J. Schneider Jr. (Ed.), Engineered Materials Handbook, vol. 4, ASM  International, Novelty, OH, United States, 1991, pp. 775-786. [33] R.O. Ritchie, X.-F. Zhang, L.C. De Jonghe, On the role of grain-boundary films in optimizing the mechanical properties of silicon carbide ceramics, Mater. Res. Soc. Symp. Proc. 819 (2004) 1.1.1-1.1.12. [34] F. Monteverde, A. Bellosi, Oxidation of ZrB2 -based ceramics in dry air, J. Electrochem. Soc. 150 (2003) B552-B559. I.B. Ban’kovskaya, V.A. Zhabrev, Kinetic  analysis of  the heat  ZrB2 -SiC composites, Glass Phys. Chem. 31 (2005) 482-488. [36] F. Peng, R.F. Speyer, Oxidation resistance of fully dense ZrB2 with SiC, TaB2 , and TaSi2 additives, J. Am. Ceram. Soc. 91 (2008) 1489-1494. [37] H.A.  Schaeffer,  Structure-property  relationships  of  the  vitreous  progress in nitrogen ceramics, NATO ASI Ser. 65 (1983) 303-321. J.B.  Berkowitz-Mattuck, High-temperature  oxidation.  III.  Zirconium  hafnium diborides, J. Electrochem. Soc. 113 (1966) 914-920. [39] D. Gao, Y. Zhang, C. Xu, Y. Song, X. Shi, Oxidation kinetics of hot-pressed ZrB2 -SiC ceramic matrix composites, Ceram. Int. 39 (2013) 3113-3119. [40] F. Peng, G. Van Laningham, R.F. Speyer, Thermogravimetric analysis of the oxidation resistance of ZrB2 -SiC and ZrB2 -SiC-TaB2 -based compositions  in the 1500-1900   C range, J. Mater. Res. 26 (2011) 96-107. [41] W.W. Pultz, W. Hertl, SiO2 + SiC reaction at elevated temperatures. Part 1. Kinetics and mechanism, Trans. Faraday Soc. 62 (1966) 2499-2504. [42] L. Silvestroni, Oxidation of ZrB2 Containing Hi-Nicalon fiber in the Temperature Range of 1200-1650   C, 2013 (unpublished raw data). [43] L. Silvestroni, D. Sciti, Oxidation of ZrB2 ceramics containing sic as particles, whiskers, or short fibers, J. Am. Ceram. Soc. 94 (2011) 2796-2799. [44] S. Hampshire, M.J. Pomeroy, Silicon nitride-grain boundary oxynitride glass interfaces: deductions from glass bulk properties, Int. J. Appl. Ceram. Technol. 10 (2013) 747-755. [45] D. Sciti, L. Silvestroni, G. Saccone, D. Alfano, Effect of different sintering aids on thermo-mechanical properties and oxidation of sic fibers-reinforced ZrB2 composites, Mater. Chem. Phys. 137 (2013) 834-842. [46] P.A. Williams, R. Sakidja,  J.H. Perepezko, P. Ritt, Oxidation of ZrB2 -SiC ultrahigh temperature composites over a wide range of SiC content, J. Eur. Ceram. Soc. 32 (2012) 3875-3883. [47] M. Takeda, A. Urano, J. Sakamoto, Y. Imai, Microstructure and oxidative degradation behavior of  silicon  carbide fiber Hi-Nicalon  type S,  J. Nucl. Mater. 258-263 (2008) 1594-1599.  resistance of  [38]   state,   in  and  \\x0c']"
},{
  "_id": 139,
  "PDF": "Oxidation behavior of a pressureless sintered ZrB2–MoSi2 ceramic composite.pdf",
  "Text": "['Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  Diletta Sciti,a) Mylène Brach, and Alida Bellosi  CNR-ISTEC, Institute of Science and Technology for Ceramics, I-48018 Faenza, Italy  (Received 17 November 2004; accepted 7 January 2005)  Ultra-refractory ceramic composites of composition ZrB2 + (5 to 20) vol% MoSi2 were produced by pressureless sintering at 1830 °C under argon atmosphere. Sintering cycles and microstructural analysis point out that at least 20 vol% molybdenum disilicide is necessary for obtaining a dense material. Thereafter, the composite 80 vol% ZrB2 + 20 vol% MoSi2 was used to test the thermal stability under oxidizing environment. Oxidation tests were carried out in flowing synthetic air in a thermogravimetric analyzer from 700 to 1400 °C with exposure time of 30 h. In the low-temperature range (700-1000 °C), the oxidation of the composite resembles that of monolithic ZrB2 ceramics, for temperatures >1200 °C the silica resulting from oxidation of molybdenum disilicide seals the sample surface, preventing zirconium diboride from fast degradation.  I.  INTRODUCTION  The design and production of new materials suitable to withstand high temperatures are nowadays stimulated by the increasing demand for applications in the fields of thermal protection systems for several industrial sectors and for advanced re-entry vehicles. Among the ultrahightemperature ceramics (UHTCs), ZrB2 displays a number of unique properties, including hardness, high thermal and electrical conductivity, and chemical stability.1-4 The ma jor prob lems encoun tered w i th ZrB2-based materials concern densification and high-temperature oxidation. Sintering of ZrB2 powders to full density is a difficult task due to the high refractoriness of this phase, unless pressure-assisted sintering procedures and temperatures higher than 2000 °C are used. Recently, ZrB2 has been successfully densified by hot pressing with the addition of phases able to improve ZrB2 sinterability, such as Ni,1 Si3N4,2 and AlN,3 while monolithic ZrB2 hardly reached a final density of 90%, even at a temperature of 1900 °C and an applied pressure of 30 MPa. On the other hand, the use of sintering aids often results in the formation of secondary phases, which can affect the oxidation resistance of the final composite. Monolithic ZrB2 oxidizes to zirconia and liquid boria following a parabolic behavior in the range of 900-1500 °C.5,6 However, at temperatures  a)Address all correspondence to this author. e-mail: dile@istec.cnr.it DOI: 10.1557/JMR.2005.0111  above 1100 °C, B2O3 starts to vaporize noticeably, and the ZrO2 porous layer is no longer protective. To overcome this problem, it has been shown that the addition of SiC is highly beneficial for improving ZrB2 oxidation resistance at temperatures >1200 °C, thanks to the formation of a protective borosilicate glass.7-12 However, the final densities of pressureless sintered ZrB2-SiC composites do not exceed 75-80%; the production of materials containing low amounts of residual porosity and controlled microstructure need pressure-assisted sintering technologies. On the basis of the above-reported considerations, finding sintering agents that allow full densification by pressureless sintering while also improving the oxidation resistance of ZrB2, would result in a great advance from a technological point of view. In this work, ZrB2 is combined with MoSi2 to investigate its suitability either as sintering aid or as a supplier of “protective” silica during the oxidation process at high temperature. Actually, MoSi2 has a combination of properties suitable for several high-temperature applications.13 As a result of its thermochemical instability, it is known to form a surface layer of silica. The silica acts as protective barrier against high-temperature oxidation, while at low temperatures, namely 500-800 °C, MoSi2 experiences pesting oxidation.14-21 Pesting disintegration can be suppressed by the addition of Al cations,17 and recently, the addition of boron to MoSi2 was shown to improve its oxidation resistance at 500 °C.18 Hot-pressed ZrB2-MoSi2 composites have been studied and it was found that the addition of 10-30 vol%  922  J. Mater. Res., Vol. 20, No. 4, Apr 2005  © 2005 Materials Research Society  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  MoSi2 improves both mechanical properties22 and oxidation resistance22-24 compared to monolithic ZrB2 materials. However, to the authors’ knowledge, no studies are reported on similar composites produced by pressureless sintering. In the present work, the results of the pressureless sintering cycles carried out on ZrB2-(5-20) vol% MoSi2 composites are presented. The best material in terms of density and microstructural features is studied, and the possible densification mechanisms are investigated. The same material has been selected to study the short-term oxidation behavior in air in the temperature range 700- 1400 °C.  II. EXPERIMENTAL PROCEDURE  A. Material  Commercial powders were used for the preparation of composite materials: ZrB2 Grade B (H.C. Starck, Karlsruhe, Germany), grain size in the range of 0.1-8 \\u242em, 1 wt% oxygen content, MoSi2 (Aldrich, Milwaukee, WI) with a mean particle size 2.8 \\u242em, and oxygen content of about 1%. The compositions prepared are reported in Table I. The powder mixture was milled for 24 h in absolute ethanol using zirconia balls. The slurry was dried in a rotary evaporator and sieved at 250 \\u242em. Pellets were prepared by uniaxial pressing followed by cold isostatic pressing under 350 MPa. Then they were pressureless sintered in a graphite furnace with atmosphere of flowing argon at different temperatures and times (Table I) to optimize the sintering cycle. Sintered densities were measured using Archimedes’ method.  B. Oxidation tests  On the basis of the sintering tests and of the microstructure of the dense materials, the composite 80 vol% ZrB2 + 20 vol% MoSi2 was selected to produce the samples for the oxidation tests. Rectangular plates sized 9.0 × 8.0 × 1.0 mm3 were cut from the sintered pellet. The surface was mechanically ground, and the measured surface roughness was about 1 \\u242em. The specimens were c leaned in u l t rason ica ted ace tone ba th , d r ied , and weighed (accuracy 0.01 mg). The oxidation tests were  TABLE I . Compo s i t ion , composites.  s in t e r ing  cy c l e ,  and  d en s i ty  o f  th e  Sample  Composition  Sintering cycle  A  B  C  95 vol% ZrB2 + 5 vol% MoSi2 85 vol% ZrB2 + 15 vol% MoSi2 80 vol% ZrB2 + 20 vol% MoSi2  1800 °C/30 min  1800 °C/60 min  1830 °C/30 min  Final  density  (%)  86.3  97.6  99.7  carried out in a thermogravimetric analyser (TG; model STA449, NETSCH, Geraetebau GmbH, Selb, Germany) in synthetic air (composition: 80 vol% N2 + 20 vol% O2, with 30 ml/min gas flow) between 700 and 1400 °C with isothermal exposure time of 30 h for each experiment, heating rate of 30 °C/min, and free cooling. The fast heating-up stage prior to the isothermal period was applied to minimize oxidation effects before reaching the target temperatures. The mass variation was recorded continuously with 10−3 mg sensitivity. The TG measurement evaluation was performed with the subtraction of buoyancy effect corrections. Oxidized and as-sintered sample surfaces were analyzed by x-ray diffraction (Cu K␣ radiation, Miniflex Rigaku, Tokyo, Japan). Sample surfaces and polished cross sections were analyzed by scanning electron microscopy (SEM; Leica Cambridge S360, Cambridge, UK) and energy dispersive microanalysis (EDS; Model INCA energy 300; Oxford Instruments, High Wycombe, UK).  III. RESULTS AND DISCUSSION  A. Microstructure of  the as-sintered materials  The amounts of 5 and 15 vol% MoSi2 were not enough for the densification of the composite since the final densities were rather low, less than 98% theoretical, calculated on the basis of the starting nominal compositions (see Table I). A final relative density of 99.7% was obtained at 1830 °C/30 min for the compositions containing 20 vol% MoSi2, which was therefore selected for microstructural characterization and oxidation tests. The x-ray diffraction pattern for this dense composite showed the presence of crystalline ZrB2 and MoSi2 and traces of Mo4.8Si3C0.6 phase usually found in MoSi2 samples containing SiC or C.25,26 The sample was polished with diamond paste to 0.25 \\u242em and then further polished with colloidal silica. This polishing procedure produces an etching effect, as shown in the backscattered electron image of Fig. 1. The globular grains are ZrB2 phase (which appears to undergo preferential removal by polishing), while the phase appearing in relief is MoSi2 (compare the EDS spectra reported). The different contrast effect exhibited by ZrB2 grains is due to their different crystallographic orientation. Secondary phases were detected by EDS analysis and include Mo-rich phases like Mo-B, Mo-Si-B, Mo-Si-C phases with different composition and Zr(C,O) phases (see the EDS spectra in Fig. 1). The estimated volume amounts of Mo-rich phases and Zr-C phases are about 3 vol% each, as calculated from image analysis (Image Pro-plus 4.0, Media Cybernetics, Silver Springs, MD). Sintering tests revealed that, indeed, MoSi2 particles improve the sinterability of ZrB2, for amounts higher than 15 vol%. Microstructural observations suggest that pressureless sintering is enhanced by the formation of  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  J. Mater. Res., Vol. 20, No. 4, Apr 2005  923  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  FIG. 1. Polished surface of pressureless sintered 80 vol% ZrB2 + 20 vol% MoSi2 composite and EDS spectra of selected areas.  liquid phase. The observed secondary phases (either Mo-B, Mo-Si-B, or Zr-C) present morphological shapes with dihedral angles values low up to 20-30°, which indicates that they (or the phases that gave rise to them) were liquid at the processing temperature, and their wettability toward the solid ZrB2 or MoSi2 grains was very high. In the system Zr-B, solid ZrB2 coexists with liquidus at temperatures higher than 1660 ± 15 °C.27 The Mo-B system is very complex, and liquid phases are likely to occur at the processing temperature.28 Even the Zr-C phase diagram shows the presence of liquid phases at temperatures of 1835 ± 20 °C.29 These phases originate from reactions between the species constituting the  starting powders. Pure MoSi2 and ZrB2 are not predicted to react even at the sintering temperature, but in the presence of carbon, and oxygen, their reaction is strongly favored. According to thermodynamic calculations,30 one possible reaction is  ZrB2共s兲 + 2MoSi2共s兲 + C共s兲 + 4O2共g兲 = ZrC共s,l兲 + 2MoB共s,l兲 + 4SiO2共s,l兲 ,  (1)  where the ⌬G is negative in the sintering temperature range (⌬G at 1500 and 1830 °C are −2264 and −2044 kJ, respectively). In the present  the sintering environment  investigation,  924  J. Mater. Res., Vol. 20, No. 4, Apr 2005  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  by a dissolution/reprecipitation mechanism of ZrB2. A liquid-phase sintering mechanism was in fact hypothesized for hot-pressed ZrB2 materials added with either Si3N4 or Si3N4/SiC,2 but in that case, secondary amorphous phases were observed on sintered samples. In the present case, microstructural observation does not permit discrimination between liquid-phase or solid-state sintering mechanisms.  B. Oxidation curves  TG curves of the composite 80 vol% ZrB2 + 20 vol% MoSi2 recorded during the isothermal run at different temperatures are shown in Fig. 3(a). The apparent weight gain could be underestimated due to the loss of volatile  is carbon rich, due to graphite heating elements crucibles and supports while oxygen impurities, detected by EDS analysis, can derive from the starting powders. Actually, more complex reactions occurred, leading to the formation of Mo-Si-B phases with variable stoichiometry. The liquid phase could also be silica derived from MoSi2 particles or borosilicate phases due to the reaction of silica with boria present on the surface of ZrB2 particles, even if in the sintered samples, evidence of either silica or borosilicate glass phases was not found. Silica pockets (as well as Mo-B and Zr-C phases) were found in hot-pressed ZrB2-20 vol% MoSi2 composites, as shown in Fig. 2. This supports the hypothesis that silicacontaining liquid phases should play a key role in the sintering of this system especially at the beginning of densification, i.e., in the 1400-1500 °C temperature range (as observed in the densification curves detected during hot pressing tests of the same powder mixtures). Then the sintering conditions (temperature and atmosphere) must have favored the vaporization and/or carboreduction of the oxide species. The sintering mechanism for ZrB2 particles, whether it be solid-state or liquidphase sintering, is not clear. It can be hypothesized that the liquid phase formation (Mo-rich phases and/or silicate glass) leads to a partial removal of boria from the surface of ZrB2 particles increasing their reactivity, and therefore, neighboring ZrB2 particles can sinter by solidstate sintering. This mechanism was already hypothesized3 for the densification of ZrB2 with addition of AlN. Alternatively, the liquid phase can favor sintering  FIG. 2. Polished surface of hot-pressed 80 vol% ZrB2 + 20 vol% MoSi2 composite, showing silica pockets, as confirmed by EDS spectrum in the insert.  FIG. 3. (a) Thermogravimetric curves of the sample 80 vol% ZrB2 + 20 vol% MoSi2 in the range 700-1400 °C; (b) corresponding squared weight gains, showing the parabolic feature of the 1200 and 1300 °C curves after t ⳱ 300 min and the pure parabolic behavior of the 1400 °C curve. The parabolic weight constants (Kp) are calculated for cycles at 1200, 1300, and 1400 °C.  J. Mater. Res., Vol. 20, No. 4, Apr 2005  925  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  species, as discussed below. The weight gain after oxidation at 1200 °C is remarkably lower than that after oxidation at 1000 °C. The kinetics of oxidation deviate from parabolic kinetics [see the squared weight changes in Fig. 3(b)], with the exception of the oxidation curve at 1400 °C satisfactorily fitted according to the equation t1/2, where w is the mass change per unit area expressed as a function of time, and Kpar is the parabolic rate constant. Figure 3(b) shows that, after a transition stage of about 300 min, the curves for oxidation at 1200 and 1300 °C are parabolic as well. The calculated parabolic weight constants are reported in the same graph.  w = Kpar  C. Oxidation product  The crystalline phases formed after oxidation are shown in Table II, along with the spectrum of the assintered material. The morphological evolution of surfaces and cross sections is reported in Figs. 4(a)-4(d) and 5(a)-5(d). According to microstructural observation, the evolution of microstructure can be divided into two regimes.  1. Samples oxidized in the temperature range 700-1000 °C  Monoclinic zirconia is the main crystalline product of oxidation. The peaks of the starting phases, namely ZrB2 and MoSi2, are still visible, but their intensity decreases with increasing oxidation temperature due to the masking effect of the growing oxide (Table II). The sample surface after oxidation at 700 °C appears covered with small crystals belonging to zirconia, but the oxide layer is porous [Fig. 4(a)]. Monoclinic ZrO2 peaks are visible in the x-ray diffraction pattern. The oxide thickness (not shown) is 1 \\u242em, confirming that the extent of oxidation at 700 °C is very limited. The surface oxide of the samples oxidized at 800 and 1000 °C [an example is reported in Fig. 4(b)] is porous and traversed by big cracks. Surface cracking is caused by several concurrent phenomena: escape of volatile species, volume expansion that follows the reaction of zirconium diboride  TABLE II. Crystalline phases of as-sintered sample and after term oxidation.  short  As-sint.  ZrB2 MoSi2 m-ZrO2 ZrSiO4 MoB  c-ZrO2 SiO2 (cristobalite)  vs  s  …  …  …  …  700  °C  s  w  vw  …  …  traces  800  °C  w  vw  s  …  …  …  1000  °C  vw  …  s  …  vw  …  …  …  traces  1200  °C  1300  °C  1400  °C  s  w  s  w  w  …  …  s  w  s  w  w  …  …  s  w  s  w  w  …  …  Legend: vw, very weak; w, weak; s, strong; vs, very strong.  to zirconium oxide, and cooling of the glassy scale. During the treatment at 1000 °C [Fig. 5(b)], a subsurface oxidation occurs: beneath the zirconia layer, large cavities are visible where new phases MoB and/or Mo-oxide formed, as found by EDS analysis. Traces of MoB peaks were detected in x-ray diffraction spectra, but due to overlap of peaks in the spectra, a clear identification of this phase is quite difficult. The cross section of the sample oxidized at 1000 °C also shows the formation of cracks well below the surface oxide [Fig. 5(b)].  2. Samples oxidized in the temperature range 1200-1400 °C.  The main feature of these specimens is that the surface starts to be covered by a continuous silica-rich glassy layer, in which small zirconia and/or zircon grains are embedded [Figs. 4(c) and 4(d)]. This layer has a thickness of less than a few micrometers and completely seals the sample surface [Figs. 5(c) and 5(d)]. Monoclinic zirconia is the main oxidation product detected in x-ray diffraction patterns; however minor quantities of MoB and zircon are also visible (Table II). Due to the thin oxide layer relative to those formed at other temperatures, the peaks of the starting phases, ZrB2 and MoSi2, are quite strong, especially at 1200 °C. Despite the presence of the surface oxide, subsurface oxidation takes place, which involves the penetration of the unreacted bulk by the glassy phase and the formation of small MoB crystals, as shown in Fig. 6. According to EDS analyses, the composition of the penetrating glass is different from the surface one, containing O, Si, B, and traces of N (the latter deriving from the oxidizing atmosphere). At higher temperatures, namely 1300 and 1400 °C, the subsurface oxidation continues down to a depth of about 100 \\u242em [Fig. 5(d)].  D. Oxidation behavior  The oxidation reactions occurring in the ZrB2-MoSi2 composite during high-temperature treatments in air induce surface and subsurface modification affected by the reactions of the two main phases constituting the composite. ZrB2 oxidizes according to the following reactions5  ZrB2共s兲 + 5 Ⲑ 2O2共g兲 = ZrO2共s兲 + B2O3共l兲  ,  B2O3共l兲 = B2O3共s兲  .  (2)  (3)  Boric acid has a low melting point (450 °C) and a high vapor pressure, which makes it vaporize at relatively low temperatures. At a temperature of around 1100 °C, liquid boria starts to vaporize noticeably.5 The oxidized surface is composed of a nonprotective zirconia-based layer. On the other hand MoSi2 oxidizes according to16 MoSi2共s兲 + 7 Ⲑ 2 O2共g兲 = SiO2共s兲 + MoO3共s兲  (4)  .  926  J. Mater. Res., Vol. 20, No. 4, Apr 2005  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  FIG. 4. SEM micrographs of the sample surfaces after oxidation for 30 h in synthetic air at selected temperatures.  FIG. 5. SEM micrographs of the cross-sections of the sample oxidized for 30 h in synthetic air at selected temperatures.  At temperatures >1000 °C, MoSi2 has excellent oxidation resistance because of a protective silica surface layer. At lower temperatures (400-800 °C), the growth of  a protective silica is not very fast and the concurrent formation/volatilization of MoO3 gives rise to very high internal stresses, which cause an enhanced oxidation of  J. Mater. Res., Vol. 20, No. 4, Apr 2005  927  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  FIG. 6. SEM micrographs of the cross section of the sample oxidized for 30 h in synthetic air at 1200 °C showing the penetration of glassy phase toward the unreacted bulk, formation of MoB crystals, and EDS of selected areas. The presence of carbon peak in the spectra is due to graphite coating.  MoSi2. The phenomenon can be so strong to induce catastrophics disintegration of the material. In the composite of the present work, according to microstructural observations, the oxidation can be divided into two regimes.  1. 800-1000 °C  In the temperature range 800-1000 °C, the formation of protective silica from reaction (4) is not fast enough to inhibit rapid oxidation. Therefore, the process is dominated by the oxidation of ZrB2 into zirconia according to reactions (2) and (3), concurrent with the formation and volatilization of MoO3 by reaction (4). B2O3, liquid at these temperatures, is known to have a protective effect, while ZrO2 is semi-protective due to its anion deficiency.9 This is confirmed by the kinetic curves, which resemble those reported for the oxidation of pure ZrB2.5 The volume expansion due to the formation of zirconia and to the oxidation of MoSi2 to Mo-oxides, as well as the escape of gaseous species, leaves voids and cracks in  the oxide layer. During oxidation at 1000 °C, the presence of newly formed crystalline MoB indicates that, besides reactions (2)-(4), the following reaction is likely to occur:  ZrB2共s兲 + 2MoSi2共s兲 + 5O2共g兲 = ZrO2共s兲 + 2MoB共s兲 + 4SiO2共s兲 .  (5)  According to thermodynamic calculations,29 reaction (5) is favored at all oxidation temperatures, with ⌬G at 700 and 1400 °C being −3542 and −2939 kJ respectively. In the temperature range 700-1000 °C, the formation of silica is very slow; therefore, especially at 1000 °C, some vaporization of molybdenum oxide14-18 and liquid boria is expected to occur.  2. 1200-1400 °C  It was previously observed that, in monolithic ZrB2 materials at 1200 °C, the evaporation of liquid boria leads to a rapid degradation.5 In the composite of the present study, however, the production of silica resulting  928  J. Mater. Res., Vol. 20, No. 4, Apr 2005  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  TABLE III. Comparison of data relative to the oxidation behavior of the composite studied in the present work (oxidized in synthetic air for 30 h) and other ZrB2-based materials reported in literature. Weight gain data for ZrB2-20 vol% MoSi2 material are obtained from the TG curves in Fig. 3(a).  Composition of various  ZrB2-based materials  ZrB2 + 4Ni1 (wt%) 55ZrB2 + 41TiB2 + 4Ni1 (wt%) ZrB2 + 20 volSiC11 (vol%) ZrB2 + 20 volSiC11 (vol%) ZrB2 ZrB2  11  11  Oxidation  cycles  (°C/h)  1000/30  1000/30  1200/2  1400/2  1200/2  1400/2  Furnace/atmosphere  Laboratory kiln  Laboratory kiln  Thermogravimetric/air  Thermogravimetric/air  Thermogravimetric/air  Thermogravimetric/air  Apparent  weight  gain (mg/cm2)  7  20  5  5  100  180  ZrB2 + 20 MoSi2 weight gain (mg/cm2)  ZrB2 + 20 MoSi2 behaviora  1.8  1.8  0.7  1.8  0.7  1.8  B  B  B/C  B/C  B  B  aB, better; C, comparable.  from oxidation of MoSi2 is fast enough to seal the sample surface and to hinder oxygen transport from the atmosphere to the reaction interface. At the beginning of the oxidation stage at 1200 °C, oxidation of ZrB2, crack formation, and evaporation of liquid boria take place. However, after about 100 min, a continuous silica-rich layer forms, which exerts a protective action and causes a drop in the weight gain. In spite of the decelerated kinetics, a subsurface reaction takes place. The generation of cracks, probably related to the formation of zirconia and Mo-oxides, favors the penetration of oxygencontaining glasses and, consequently, the contact of oxygen with zirconium diboride and molybdenum disilicide particles also in the unreacted bulk. At the interface between the oxide and the unreacted bulk, the partial pressure of oxygen is rather low, which leads to selective oxidation of Si; therefore, MoSi2 is depleted of Si, and Mo reacts with the boron of ZrB2 to form MoB. Unlike the case of ZrB2-SiC composites,9 little borosilicate glass was observed in the oxide layer. In contrast, silicarich regions associated with the presence of MoB crystals were detected. This suggests that the formation of MoB through reaction (5) is favored over the formation of borosilicate glass. Meanwhile, at the surface of the sample where zirconia crystals are embedded in a continuous and partially amorphous silica-rich layer, further reaction can occur among the oxidation products to form crystalline zircon. According to previous studies,31 when zirconia grains are embedded in a silica layer, interstitial silicon diffuses and dissolves into crystalline zirconia until the solubility limit is reached; thereafter zircon precipitates. The resulting oxide consists of zirconia grains, surrounded by a ZrSiO2 shell, which, in turn, is embedded in the amorphous silica. The possible reaction is  ZrO2  + SiO2 共amorphous兲 = ZrSiO4  .  (6)  The nearly parabolic kinetics of oxidation at 1200 and 1300 °C and the pure parabolic kinetics found at 1400 °C indicate that the dominant mechanism may be diffusion  of oxygen through the silica-based oxide layer, as observed for ZrB2-SiC composites. A comparison of the oxidation behavior between the composite of the present work [on the basis of the TG curves in Fig. 3(a)] and other ZrB2-based materials is given in Table III. With reference to monolithic ZrB2, in the work of Monteverde et al.,1 no oxidation rates are reported for Ni-doped materials. However, weight gain curves and microstructural observations indicate that during oxidation at 1000 °C/30 h, the resistance of the ZrB2-20 vol% MoSi2 composite is higher, although the oxidizing atmosphere is different. If the weight gain achieved in similar conditions of time/temperature is compared for the two composites ZrB2-20 vol% MoSi2 and ZrB2-20 vol% SiC, their oxidation resistance in the 1200-1400 °C temperature range with a 2 h isothermal stage11 is similar. These are only indications, since all the data collected from literature concern oxidation tests carried out in different experimental conditions (in term of temperature, oxidizing atmosphere, and exposure time) and, besides, in some cases, the microstructure of the materials (i.e., porosity and phase composition) is not completely defined.  IV. CONCLUSIONS  ZrB2-based composites were successfully densified without pressure at 1800 °C, with addition of 20 vol% of molybdenum disilicide. Sintering is hypothesized to be aided by the formation of liquid phases, either as borosilicatic phase (due to B2O3 in ZrB2 and SiO2 in MoSi2) or Mo-rich phases due to reaction between the starting ZrB2 and MoSi2. Oxidation tests proved that the composite 80 vol% ZrB2 + 20 vol% MoSi2 can withstand temperatures up to 1400 °C under an oxygen-containing environment. At lower temperatures, the composite oxidation behavior resembles that of monolithic ZrB2, since zirconium oxide is the main oxidation product. At temperatures >1200 °C, due to the oxidation of MoSi2, appreciable amounts of  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  J. Mater. Res., Vol. 20, No. 4, Apr 2005  929  \\x0c', 'D. Sciti et al.: Oxidation behavior of a pressureless sintered ZrB2-MoSi2 ceramic composite  glass form, which improves the oxidation resistance of the material. At higher temperatures, a concern is the subsurface oxidation, involving glass penetration and formation of the MoB phase. This study points out that this ZrB2-20 vol% MoSi2 composite has the advantage of sintering at relatively low temperature compared to monolithic ZrB2 or ZrB2-SiC materials. Furthermore, the oxidation resistance of the composite containing MoSi2 is higher than that of monolithic ZrB2 and comparable to that of hot pressed SiCreinforced ZrB2 materials.  ACKNOWLEDGMENTS  The work was supported by the European Project Research Training Network HPRN-CT-2000-00044 “Composite Corrosion.” The research contract of M. Brach was funded by the same project. The authors wish to thank Andrea Balbo for the TG measurements and Frederic Monteverde for useful discussion on microstructure.  REFERENCES  1. F. Monteverde, A. Bellosi, and S. Guicciardi: Processing and properties of zirconium diboride-based composites. J. Eur. Ceram. Soc. 22, 279 (2002). 2. F. Monteverde, S. Guicciardi, and A. Bellosi: Advances in microstructure and mechanical properties of zirconium diboride based ceramics. Mater. Sci. Eng. A346, 310 (2003). 3. F. Monteverde and A. Bellosi: Beneficial effects of AlN as sintering aid on microstructure and mechanical properties of hotpressed ZrB2. Adv. Eng. 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Jessen (ACERS, 22, Westerville, OH, 2001), p. 221-228. 11. M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykosky,  Engineering and Science Proceedings,  in Ceramic  and  S.J. Causey: Mechanical thermal and oxidation properties of refractory hafnium and zirconium compounds. J. Eur. Ceram. Soc. 19, 2405 (1999). 12. C.R. Wang, J.M. Yang, and W. Hoffman: Thermal stability of refractory carbide/boride composites. Mater. Chem. Phys. 74, 272 (2002). 13. Y-L. Jang and E.J. Lavernia: Review: Processing of molybdenum disilicide. J. Mater. Sci. 29, 2557 (1994). 14. C.G. McKamey, P.F. Tortorelli, J.H. Devan, and C.A. Carmichael: A study of pest oxidation in polycrystalline Mo Si2. J. Mater. Res. 7, 2747 (1992). 15. P.J. Meschter: Low-temperature oxidation of molybdenum disilicide. Metall. Trans. 23, 1763 (1992). 16. T.C. Chou and T.G. Nieh: Mechanism of MoSi2 pest during low temperature oxidation. J. Mater. Res. 8, 214 (1993). 17. K. Natesan and S.C. Deevi: Oxidation behaviour of molybdenum silicides and their composites. 18. Y.T. Zhu, L. Shu, and D.P. Butt: Kinetics and products of molybdenum disilicide powder oxidation. J. Am. Ceram. Soc. 85, 507 (2002). 19. K. Kurokawa, H. Houzumi, I. Saeki, and H. Takahashi: Low temperature oxidation of fully dense and porous MoSi2. Mater. Sci. Eng. A261, 292 (1999). 20. K. Yanagihara, T. Maruyama, and N. Kazuhiro: Effect of third elements on the pesting suppression of Mo-Si-X intermetallics (X ⳱ Al, Ta, Ti, Zr and Y). 21. H.M. Yokota, T. Kudoh, and T. Suzuki: Oxidation resistance of boronized MoSi2. Surf. Coat. Technol. 169-170, 171 (2003). 22. A.L. Chamberlain, W.G. Fahrenholtz, and G.H. Hilmas: Characterization of zirconium diboride-molybdenum disilicide ceramics,  Intermetallics 4, S133 (1996).  Intermetallics 8, 1147 (2000).  in Ceramic Matrix Composites, Ceramic Transac in Advances tions ,  26.  by N .P . Bansal , J .P . Singh , W .M . Kriven , and (Am. Ceram. Soc., 153, Westerville, OH, 2003),  edited H. Schneider p.299. 23. P.T.B. Shaffer: An oxidation resistant boride composition. Ceram. Bull. 41, 96 (1962). 24. M.M. Opeka, I.G. Talmy, and J.A. Zaykoski: Oxidation-based materials selection for 2000°C + hypersonic aerosurfaces: theoretical considerations and historical experience. J. Mater. Sci. 39, 5887 (2004). 25 . D . Sc i t i , S . Gu icc iard i , C . Me landr i , and A . Be l los i : H ightemperature resistant in the AlN-SiC-MoSi2 system. J. Am. Ceram. Soc. 86, 1720 (2003). J.I. Lee, N.L. Hecht, and T-I. Mah: In-situ processing and properties of SiC/MoSi2 nanocomposites. J. Am. Ceram. Soc. 81, 421 (1998). 27. A.E. McHale: Phase Diagram for Ceramists (Am. Ceram. Soc., 10, Westerville, OH, 1994), p. 144. 28. A.E. McHale: Phase Diagram for Ceramists (Am. Ceram. Soc., 10, Westerville, OH, 1994), p. 136. 29. A.E. McHale: Phase Diagram for Ceramists (Am. Ceram. Soc., 10, Westerville, OH, 1994), p. 174. 30. HSC Chemistry for Windows 5, Outokumpu Research Oy, Pori, Finland (2002). 31. C. Veytizou, J.F. Quinson, O. Valfort, and G. Thomas: Zircon formation from amorphous silica and tetragonal zirconia: Kinetic study and modelling. Solid State Ionics 139, 315 (2001).  930  J. Mater. Res., Vol. 20, No. 4, Apr 2005  http://journals.cambridge.org  Downloaded: 01 Mar 2015  IP address: 138.251.14.35  \\x0c']"
},{
  "_id": 140,
  "PDF": "Oxidation behavior of graphene nanoplatelet reinforced tantalum carbide composites in high temperature plasma flow.pdf",
  "Text": "['C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  A v a i l a b l e a t w w w . s c i e n c e d i r e c t . c o m  ScienceDirect  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c a r b o n  Oxidation behavior of graphene nanoplatelet  reinforced tantalum carbide composites in high  temperature plasma flow  Andy Nieto a, Amit Kumar b, Debrupa Lahiri a, Cheng Zhang a, Sudipta Seal b, Arvind Agarwal a,*  a Plasma Forming Laboratory, Nanomechanics and Nanotribology Laboratory, Mechanical and Materials Engineering, 10555 West Flagler  Street, EC 3464, Florida International University, Miami, FL 33174, USA  b AMPAC and Nanoscience Technology Center, 4000 Central Fl Boulevard, University of Central Florida, Orlando, FL 33816, USA  A R T I C L E  I N F O  A B S T R A C T  Article history:  Received 17 August 2013  Accepted 8 October 2013  Graphene nanoplatelets (GNP) reinforced tantalum carbide (TaC) composites are exposed to  a high temperature plasma flow in order to evaluate the effects of GNP on the oxidation  behavior of TaC at conditions approaching those of hypersonic flight environments. The  Available online 19 October 2013  addition of GNP is found to suppress the formation of the oxide layer by up to 60%. The high  thermal conductivity of GNPs dissipates heat throughout the sample thereby reducing ther mal gradients and reducing the intensity of heating at the surface exposed to plasma.  In  addition, GNPs enhance oxidation resistance by providing toughening which suppresses  crack formation and bursting that accelerates oxidation. Scanning electron microscopy  (SEM) and high resolution transmission electron microscopy (HR-TEM)  reveal  that GNPs  have the ability to survive the intense high temperature of the plasma. GNPs are believed  to seal oxide grain boundaries and hinder the further influx of oxygen. GNPs also provide  nano sized carbon needed to induce the localized reduction of Ta2O5 to TaC. Micro com puted X-ray tomography (MicroCT) validates that the above mechanisms protect the under lying unoxidized material from the structural damage caused by thermal shocks and high  shear forces, by reducing thermal gradients and providing toughness.  Ó 2013 Elsevier Ltd. All rights reserved.  1.  Introduction  points. The high density of TaC (14.65 g/cm3) [7] makes it more  promising for leading edge components as compared to lower  Tantalum carbide (TaC)  is an ultra high temperature ceramic  (UHTC) with potential applications in high temperature aero density UHTCs  such as ZrB2 and HfB2 11.21 g/cm3 [8] respectively) as it facilitates the shifting of the  (6.12 g/cm3  [8]  and  space systems such as hypersonic missiles, planetary entry  aerodynamic center forward. The extreme shock waves experi vehicles, and SCRAM jet engines [1-3]. The high melting point (3880 °C) [4] and resistance to chemical attack make TaC an ideal  candidate for surviving the extreme conditions of hypersonic flows. TaC oxidizes at temperatures as low as 400 °C (Ta2O5), however, both the oxides (1872 °C [5]) and other stoichiometric variants of tantalum carbide (>2500 °C [6]) possess high melting  enced during hypersonic flight mandate that  the candidate  material should possess excellent mechanical properties along side resistance to oxidation and chemical attack. TaC has high  hardness  (13.5-20 GPa)  [9]  and high elastic modulus  (477-  560 GPa) [9], however the low fracture toughness (4-5 MPa m1/2)  [10] makes it currently unviable for aerospace applications.  * Corresponding author: Fax: +1 305 348 1932.  E-mail address: agarwala@fiu.edu (A. Agarwal). 0008-6223/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.carbon.2013.10.010  \\x0c', 'Previous  studies  in our  lab reported the use of  carbon  nanotubes (CNT) as reinforcement for TaC [7,11]. It was found  that short CNTs that transformed into graphene platelets re sulted in the largest improvements in toughness and rupture  strength. Graphene nanoplatelets are carbon allotropes con sisting of few (20-40) layers of graphene with a high aspect ra tio  [12]. Graphene nanoplatelets have  been shown to  be  promising  fillers  for  improving  the  fracture  toughness  of  Si3N4 [13-17], Al2O3 [18-26], and ZrB2 [27] matrix composites.  In our previous work [28,29] Graphene nanoplatelets are used  to reinforce TaC in order  to improve  toughness, damping  behavior, and consolidation. Addition of GNP leads  to in creased densification,  reduced grain growth, and enhance ments  in toughness  and damping of up to  99% [28]  and  300% [30] respectively. The high electrical and thermal con ductivities of GNPs aid in uniformly sintering the composites  without  inducing  phase  transformations. GNP wraps  the  grains and inhibits grain growth. The primary mechanisms  responsible  for  improving  the  toughness  and damping of  the TaC-GNP composites are the unique energy dissipating  mechanisms  of GNP  such as  sheet  bending,  sliding,  and  pull-out.  While improving the mechanical properties of TaC is a cru cial step in making it a viable material, it is not sufficient if the  oxidation properties are not also improved upon. The improve ment of the oxidation behavior is crucial as formation of oxi des  consume  the high strength parent material while  the  oxides themselves have low structural integrity. When evaluat ing the oxidation behavior of a UHTC it is necessary to go be yond the conventional  isothermal furnace exposure tests and  evaluate the materials response in an environment approach ing  the  conditions  experienced  during  hypersonic  flight.  UHTCs in service will  likely fail by a combination of melting,  ablation, spalling, sublimation, and oxidation [30]. Several oxi dation tests have been carried out on UHTCs under extreme  aerothermal  conditions using  oxyacetylene  torches  [31-34],  plasma arc jets [35-39],  subsonic plasma wind tunnels [40],  and supersonic plasma wind tunnels [30,41-45]. These tests  closely approximate the intense aero-thermal  loads experi enced during hypersonic flight because of the extreme temper atures and high speeds of  the flows. The high speeds of  the  flow during these tests provide high shear stresses that can  influence the microstructural evolution and lead to mechanical  failure. The temperatures of the flow also typically exceed 2000 °C which is high enough to melt most UHTC oxides.  While only one  aerothermal oxidation study has  been  done on TaC to date [34], several studies have been conducted  on ZrB2-SiC composites [32,37,40-42,44] and hafnium based  UHTCs [31,38]. Studies comparing oxidation of ZrB2-SiC with  Si3N4-SiC in a high temperature supersonic plasma flow re vealed how the phase transformations and superior thermal  conductivity of ZrB2-SiC resulted in superior oxidation behav ior [40]. The enhanced thermal conductivity leads to less se vere  thermal  gradients while  the  formation of  boria  and  silicates provides protective glassy phases  that  inhibit  the  penetration of oxygen.  In this work, TaC-GNP composites are exposed to a high  temperature plasma flow for 60 s. The plasma is generated  by a direct current (DC) plasma spray gun and reaches temperatures of \\x182500 °C. Oxidation behavior  is  characterized  in terms of oxide layer formation, oxide layer microstructure,  and thermal gradients during testing. These high tempera ture oxidation tests are complimented by steady state isothermal TGA experiments done for 120 min at 1000 °C. It is found  that GNPs improve the oxidation behavior by providing struc tural  toughness,  enhancing  the  thermal  conductivity,  and  inducing localized reducing conditions.  2.  Materials and methods  2.1.  Materials  The TaC powder (1:1 Ta:C ratio) used was obtained from Infra mat Advanced Materials LLC, CN, USA and had a purity of  99.7% (major impurities were <0.15% free carbon, 0.15-0.30%  oxygen,  and  <0.3%  niobium).  Graphene  nanoplatelets  (xGNP-M-5) were obtained from XG Sciences, Lansing, MI,  USA. The GNP particles have an average diameter of 15 lm  and a thickness of 6-14 nm (20-40 graphene layers). In accor dance with  the  recent  editorial  on  the  nomenclature  of  graphenic materials  [46],  it  is  clarified here that  the term  GNP should indeed refer to graphite nanoplatelets rather than  graphene nanoplatelets. The GNPs here consist of over 10  graphene layers and therefore do not merit the term graph ene as per the editorial. However, this work retains the term  graphene for consistency with our previous works. The GNPs  have a relative surface area of 120-150 m2/g [12] and an aspect ratio of \\x181500. The GNPs have functional groups at the edges  consisting of approximately 0.075% ether, 0.035% carboxyl,  and 0.03% hydroxyl groups [12]. TaC reinforced with 1 vol.%  (TaC-1G), 3 vol.% (TaC-3G), and 5 vol.% (TaC-5G) of GNP are  synthesized by mixing TaC powder and GNP powder by wet  chemistry methods. The composite powders are consolidated  to  bulk  samples  by  spark plasma  sintering  (SPS) using a 1850 °C,  20 mm graphite die with sintering parameters of  80-100 MPa, and a holding time of 10 min.  2.2.  Oxidation tests  2.2.1.  High temperature plasma flow tests  The high temperature plasma jet tests are conducted using a  SG 100 DC plasma gun (Praxair, Danbury, CT, USA) operating  at a plasma power of 32 kW. Plasma is generated with the pri mary gas, argon, flowing at 56 slpm and the secondary gas,  helium, flowing at 60 slpm. This plasma gun is typically used  to deposit thermal sprayed coatings. Plasma spray guns gentemperatures above 10,000 °C and deposit  erate plasma at  powders  at  speeds  of  75-300 m/s  [47], making  them an  excellent  choice  for  simulating  the  extreme  conditions  experienced during hypersonic flight. The setup for  testing  the TaC-GNP  composites  is  shown in Fig.  1. The plasma  gun is set at a stand-off distance of 50 mm relative to the  TaC-GNP sample. This stand-off distance is sufficiently large  that it ensures the sample will not be fully immersed in the  plasma flow and will  therefore be  exposed to substantial  amount of oxygen from the atmosphere. The TaC-GNP sam ple is held in position using a tungsten fixture and exposed  for 60 s.  The  tungsten fixture has  a  small  1 mm diameter hold  drilled  into  the  backside  in  order  to  accommodate  a  C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  399  \\x0c', 'thermocouple for measuring the backside temperature in situ.  A K-type thermocouple  (KMQSS-02OU, Omega Engineering  Inc., Stamford, CN, USA)  is inserted into the hole and bent  at an upwards angle in order to ensure proper contact with  the sample backside. Thermal tape is used in order to protect  the thermocouple from thermal damage. The backside temperatures do not exceed the thermocouple’s limit (\\x181200 °C)  during the 60 s oxidation test; however the temperatures on  the front side are significantly higher and therefore unsuit able for measurement via thermocouple. The front side of  the sample is not directly measured in this experimental set up.  Instead,  the front side temperature is gauged by using  accuraspray in-flight particle diagnostic sensor (Tecnar Auto mation Ltd, QC, Canada) to estimate the temperature at a dis tance of 50 mm from the plasma gun. Accuraspray temperature is measured to be 2575 ± 55 °C. Assuming a stea dy state condition is reached after some time at the front sur2575 ± 55 °C can be  face,  approximated  as  the  front  side  temperature in order to qualitatively gauge the temperature  gradient. The desire to evaluate as large a test sample as pos sible combined with the limited amount of material available  from the TaC-GNP compacts dictates  the irregular  sample  shape shown in Fig. 1c. The samples are quarter-ellipses with lengths of \\x1815 mm and minor axis major axis lengths of \\x1810 mm. The samples have a thickness of \\x181.0-1.5 mm.  2.2.2.  Steady state  isothermal  thermal gravimetric analysis  (TGA)  Thermal gravimetric analysis (TGA) was done using an SDT  Q600 (TA Instruments, New Castle, DE, USA). The samples were  placed in an alumina crucible during the tests. The temperature is rapidly increased (heating rate >250 °C/min) to 1000 °C  and held at steady state equilibrium. The initial weight is re corded once the steady state equilibrium is reached. Samples used for TGA testing has an initial starting mass of \\x1810.0-  15.0 g, samples were obtained by fracturing the samples and  obtaining small chips.  It is emphasized that the TGA study is  not  the central  focus of  this work. TGA studies are used to  compliment  the plasma flow oxidation in order  to further  strengthen the proposed GNP oxidation mechanisms.  2.3.  Structural characterization  Cross sections of  the TaC-GNP oxidized samples were pre pared by grinding and polishing to a 0.3 lm finish using SiC  polishing papers (up to 600 grit size) and diamond slurries.  The cross sections were examined in a Buehler Versamet 3  Optical Microscope (Illinois, USA)  in order  to measure the  thickness of  the oxide layers. Scanning electron microscopy  of the TaC-GNP samples before and after oxidation was done  using  a  JEOL  JSM-633OF field emission scanning  electron  microscope (FE-SEM) with an operating voltage of 15 kV. En ergy dispersive spectroscopy (EDS) of  the oxidized TaC-GNP  cross sections was done in a JEOL JSM 5900LV SEM equipped  with an EDAX EDS system. Phase analysis is done by X-ray  diffraction (XRD) using a Siemens D-500 X-ray diffractometer  operating  at  a  current  and  voltage  of  40 mA and  40 kV  respectively.  Transmission  electron  microscopy  of  the  oxidized TaC-GNP samples was done using a Philips/FEI Tec nai high-resolution TEM operating at a voltage of 300 keV.  Sample preparation consisted of crushing the brittle oxides  to a fine powder and dispersing into an acetone solution via  ultrasonication. Forward and inverse fast Fourier transforms  (FFT) are done using Gatan Inc. Digital Microscope software  in order to accurately measure lattice plane spacings. Micro  computed X-ray tomography (MicroCT) was done at XRADIA  (Pleasanton, California, USA) using a VersaXRM-500. MicroCT  scan was performed with a voxel size of 0.58 lm, a field of  view of 0.58 mm, and X-ray beam energy of 70 kV. Microhard ness testing of the oxide layers was done using a HXD-1000  TMC microhardness tester (Shanghai Taiming Optical Instru ment Co. Ltd, Shanghai, China) using a Vickers tip in order to  characterize the structural integrity after oxidation. A load of 100 g (\\x181 N) was used with a holding time of 15 s.  3.  Results and discussion  3.1.  Pre-oxidation microstructural characterization  The synthesis and characterization of TaC-GNP composites is  briefly described here, details can be found elsewhere [28].  The resulting TaC-GNP compacts consolidated by spark plasma sintering have high relative densities, ranging from \\x1894% for pure TaC to \\x1899% for TaC-5G. The grain sizes of the compacts ranges from \\x184.5 lm for TaC to \\x181.5 lm for TaC-3G and  TaC-5G. The GNPs form networked structures such as that  shown in Fig. 2a,  in both the TaC-3G and TaC-5G structures.  The GNPs wrap  and weave  around  grains  as  shown  in  Fig. 2b, grain wrapping inhibits grain growth and promotes  more effective sintering because of the high thermal and elec trical conductivity of GNPs. The networked GNP structures  enable GNP  energy  dissipating mechanisms  to  effectively  toughen the composite structure.  3.2.  Oxide layer formation  Fig. 3 shows TaC and TaC-5G samples post oxidation. It can be  seen that TaC and TaC-5G samples survived 60 s exposure to  the  high  temperature  plasma  flow. No  visible  structural  damage, such as chipping or spallation, has occurred. The  Fig. 1 - Setup for plasma flow oxidation experiments,  (a)  tungsten fixture for holding samples, small hole in the back  allows for insertion of thermocouple to record backside  temperature in situ,  (b) typical geometry of TaC-GNP  samples tested and (c) plasma gun at a standoff distance of  50 mm to sample.  (A color version of this figure can be  viewed online.)  400  C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  \\x0c', 'C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  401  Fig. 2 - SEM micrographs showing distribution of GNPs in TaC-GNP composites, (a) GNPs wrap and weave around TaC grains  to form a networked structure in TaC-3G and (b) GNP grain wrapping inhibits grain growth and provides toughness.  samples  are  fully  covered  by  a  rough white  oxide  layer  which is confirmed by XRD, shown in Fig. 4,  to be mostly  Ta2O5. Ta2O5 formation occurs by the simple oxidation reac tion given in Eq.  (1):  2TaC þ 9 2  O2 ! Ta2O5 þ 2CO2  ð1Þ  other minor phases present include TaO3, Ta4C3, TaC, Ta, and  free carbon (graphite). Traces of melting can be seen at  the  edges of the samples in Fig. 2, as evidenced by the presence  of a glassy phase. The glassy phase is the result of the oxide  melting and rapidly solidifying once the plasma is  turned  off. The presence of glassy phases indicates the temperature experienced is in excess of 1872 °C, the melting temperature  of Ta2O5 [5]. The primary quantification of the oxidation per formance of  the TaC-GNP composites is done by measuring  Fig. 4 - X-ray diffractions (XRD) patterns of TaC-GNP  the thickness of  the formed oxide layer. The thickness of  samples after exposure to high temperature plasma flow for  the front oxide layer is provided in Fig. 5. The front side oxide layer is of interest as it was directly exposed to the \\x182500 °C  plasma flow.  It can be seen that the GNP reinforced samples  have oxide layers that are significantly thinner than the TaC  60 s.  (A color version of this figure can be viewed online.)  localized reducing conditions, which is elucidated in follow sample. The insets show optical micrographs of the cross sec ing sections.  tions of the thickest oxide layer (TaC), and the thinnest oxide  layer (TaC-3G). The addition of GNP thus suppresses the for mation of  the oxide layer by up to 60%. The differences in  oxide layer  formation with varying GNP content are minor  and not a focus of discussion. The addition of GNP is believed  to enhance the oxidation resistance and suppress the forma tion of the oxide layer due to an enhanced thermal conductiv ity, microstructural  interactions,  and  the  formation  of  3.3.  GNP oxidation mechanisms  3.3.1.  Enhanced thermal conductivity  GNP has a thermal conductivity (5.3 · 103 W/m K)  [22] that is  two orders of magnitude higher than TaC (30 W/m K)  [48].  In  our previous work [28]  it was postulated that GNP increased  the thermal conductivity of  the composite powder  leading  Fig. 3 - TaC-GNP samples after exposure to high temperature plasma flow for 60 s, (a) TaC and (b) typical TaC-GNP samples -  TaC-5G.  (A color version of this figure can be viewed online.)  \\x0c', '402  C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  GNPs can be seen to increase the thermal conductivity, with  the effect leveling off at 3 vol.% GNP.  The effect of enhanced thermal conductivity on the oxida tion behavior is schematically illustrated in Fig. 6b.  Improved  thermal conductivity reduces the oxidation of  the TaC-GNP  by more effectively transferring heat  throughout  the sample.  The transferring of heat  is especially effective in the TaC-3G  and TaC-5G samples where GNPs were seen to form networked  structures [28] as shown in Fig. 2. As illustrated in Fig. 6b, by  effectively transferring heat away from the front of the sample,  the regions with the most intense heating is greatly reduced.  The lower intensity heating provides less energy for the oxida tion of TaC to Ta2O5. The effective conduction of heat away  Fig. 5 - Front side and back side oxide layer thickness of  from the leading edge is essential  for  the hypersonic flight  TaC-GNP samples. Insets show optical micrographs of  applications for which the use of TaC is intended for [3].  In a  thickest oxide (TaC) and thinnest oxide (TaC-3G). The scale  hypersonic body,  the extreme heating conditions are drasti bar on insets reads 200 lm. (A color version of this figure can  cally reduced away from the stagnation point, making it desir be viewed online.)  to more uniform heating during sintering process. In the pres ent study, the effects of thermal conductivity are directly ob served by analyzing the thermal gradients experienced during  the exposure to the plasma flow. As explained in Section 2.2.1,  the front side temperature was assumed to reach equilibrium  with the plasma flame and the backside temperatures were  measured in situ during exposure to the plasma flow. The  thermal gradients of  the TaC-GNP samples are presented in  Fig. 6a. The time zero is taken once the oxidation tests have  begun and hence the starting gradients are not  the same.  The small oscillations in the thermal gradient at  the onset  of testing are due to the initial turbulence of the plasma flow.  It can be seen that the largest thermal gradient is experienced  by the TaC sample while the TaC-3G and TaC-5G samples  have  the  lowest. Given that  experimental  conditions  and  sample sizes were similar,  the thermal gradient  is inversely  proportional  to  the  thermal  conductivity. The  addition of  able to transfer heat from the high flux stagnation point to the  relatively ‘cool’ components downstream of the flow.  3.3.2.  GNP microstructural  interactions  In order to understand GNP microstructural mechanisms we  first  compare  the microstructures of  the oxidized TaC and  the TaC-GNP composites. SEM micrographs of the surface di rectly exposed to the plasma flow are provided in Fig. 7. Fig  7a and c, show how the morphology of the oxide varies dra matically between TaC and GNP reinforced TaC samples. Multi ple chasms, 7-26 lm in span (avg. 10.4 lm), can be seen on the  TaC sample. High magnification images in Fig. 7b, show that  the material  is peeling outwards indicating that the chasm is  a result of a bursting phenomenon. The inset in Fig. 7b, shows  that the area inside the chasm is heavily oxidized.  It is there fore believed that  the bursting is caused by an accumulation  of CO and CO2 gases trying to escape the sample. These large  chasms accelerate the oxidation process by providing an easier  route for oxygen infiltration into the sample.  Fig. 6 - (a) Thermal gradients of TaC-GNP samples, backside temperatures were recorded in situ while front side temperatures  were assumed to reach and remain close to the flame temperature and (b) schematic illustrating the change in heat transfer  with higher thermal conductivity. Higher thermal conductivity of GNPs enables higher heat dissipation throughout the  sample, thereby decreasing the intensity of heat near the front surface. Decreased heat at the front surface reduces the  amount of energy available to drive the oxidation of TaC.  (A color version of this figure can be viewed online.)  \\x0c', 'C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  403  Fig. 7 - Microstructures of TaC-GNP oxidized samples, (a) TaC sample displays large ‘burst out’ regions, (b) area around burst  appears to have been in a molten state, inset shows region inside burst is heavily oxidized, (c) TaC-GNP samples display large  cracks with some areas bulging out and (d) regions with heavy cracking are peeling but not bursting, inset shows region near  bulge maintain crystalline structure.  In contrast, the TaC-GNP samples show no signs of large  structural  integrity. This enhanced strength allows the GNP  chasms; cracks ranging in intensity are seen throughout the  reinforced samples to better  resist bursting and spallation.  sample. Fig. 7d, provides a high magnification image of one  The lower hardness of  the TaC-5G sample as compared to  of the cracks and it can be seen that the area around the crack  the other GNP reinforced samples is attributed to the uneven  is bulging out, however it is not severe enough to expose the  distribution of GNPs which causes some areas to be devoid of  underlying oxide layer. The largest cracks where bulging oc GNPs and hence the toughening effect to be absent.  curs have widths of 1-5 lm (avg. 3.9 lm), which are signifi To further understand the role of GNPs in enhancing oxi cantly  smaller  than the deep chasms  found in TaC. The  dation resistance, SEM analysis was used to gauge whether  structure of the bulged region in the TaC-GNP samples is also  GNPs survived the harsh environment due to plasma flow.  vastly different from the edge of the burst regions in the TaC  Fig. 9 provides SEM and TEM micrographs of several carbon  sample. The edge of  the TaC bursts shows a smooth glassy  structures found in the TaC-GNP oxide layers. Fig. 9a, shows  structure indicating that partial melting has occurred as a re sult  of  the  severe  temperatures. The TaC-GNP  structures  however, as can be seen in the inset of Fig. 7d, maintain the  crystalline structure of the oxide. As discussed in the previous  section, the higher thermal conductivity of the TaC-GNP sam ples enhances heat flow away from the front surface leading  to lower temperatures near the front surface.  In addition to severe heating conditions,  the initiation of  bursting requires that the pressure of gases accumulating be neath the surface exceeds  the strength of  the oxide.  It  is  known that the TaC-GNP samples are up to 99% tougher than  the TaC samples [28]. It is possible that some of this toughen ing effect carries over in the formed oxide structure. The por ous structure makes it unviable for evaluating the toughness  by the Vickers Indentation Method. However, measuring the  microhardness of  the oxide structures gives a measure of  the structural  integrity of the formed oxide. The microhard ness results of  the oxide layers are presented in Fig. 8. The  variation is high due to the highly porous nature of the oxide;  however the trend clearly indicates that  the GNP reinforced  samples have a higher hardness and therefore a superior  Fig. 8 - Microhardness of TaC-GNP oxide layers. Higher  hardness values in GNP reinforced samples indicate higher  structural integrity. Less severe cracking was observed in  GNP reinforced sample microstructures.  (A color version of  this figure can be viewed online.)  \\x0c', '404  C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  Fig. 9 - SEM and TEM analysis showing survival of GNPs and other carbonaceous structures,  (a) SEM of smooth undamaged  GNP,  (b) SEM of charred and heavily wrinkled GNP,  (c) SEM of carbon band embedded in oxide structures and (d) TEM of fine  carbon structures among the TaC oxide particles.  (A color version of this figure can be viewed online.)  a 15 lm GNP that appears to have undergone little to no dam inhibit  the influx of oxygen through grain boundaries. This  age during the oxidation process. Fig. 9b, shows a GNP that  GNP grain sealing mechanism is most effective if GNP survive  has appeared to become charred and lightly damaged as a re the harsh conditions. The mechanism by which grain sealing  sult of  the plasma flow. Evidence of finer carbon structures  inhibits the influx of oxygen is schematically presented in  originating from the GNP are shown in the form of embedded  Fig. 10. Oxygen encountering a sealed grain boundary may  submicron carbon bands in Fig. 9c, and in the form of 100 nm  be forced to seek a more arduous path around the sealed  sized particles in the TEM micrograph in Fig. 9d. These micro grain boundary.  graphs provide substantial evidence that GNP have the ability  The oxygen may also diffuse through the GNP which re to survive to some extent the harsh conditions of the plasma  quires more  energy than diffusing  through a high energy  flow. The survival of GNPs can be partially attributed to the  insulating effect of the TaC matrix. The survival of GNPs indi cates that GNPs may be able to play a more active role in  resisting oxidation, which is elucidated below.  It is known that GNPs wrap around grains and form weav ing networked structures [13-20,28]. It is postulated here that  wrapped GNP can effectively ‘‘seal’’ the grain boundaries and  Fig. 10 - Schematic illustrating how grain sealing  Fig. 11 - TEM micrograph showing carbon structures  mechanism hinders diffusion of oxygen into the TaC-GNP  wrapping and effectively sealing two TaC oxides grains.  structure.  (A color version of this figure can be viewed  TaO3 spacing is that of (0 4 0) planes.  (A color version of this  online.)  figure can be viewed online.)  \\x0c', 'C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  405  severe and therefore more GNPs are expected to survive, this  results in the superior oxidation resistance with higher GNP  content. The initial  increased oxidation rate in the TaC-GNP  samples is due to the higher thermal conductivity of the sam ples which increases heat flow into the sample.  3.3.3.  GNP induced reducing environment  The addition of carbon to carbides has been known to sup press the formation of oxides during sintering [49-51]. Simi larly the addition of nano-scale carbon can have the effect  of suppressing the formation of oxides in TaC by inducing a  localized reducing environment. The reduction of Ta2O5  to  TaC occurs by the equation given in Eq.  (2):  Fig. 12 - Thermal gravimetric analysis (TGA) at 1000 °C  shows GNP reinforced samples oxidize to a lower extent.  GNP grain sealing is dominant mechanism at this  temperature range.  (A color version of this figure can be  viewed online.)  grain boundary. The GNP may also oxidize as well as would be  Ta2O5 þ 7C ! 2TaC þ 5CO  ð  DG ¼ 1142:3 \\x00 0:823 T  Þ;  ð  T > 1135 \\x0eC  Þ ð2Þ  The presence of a high temperature plasma flow in air  makes reduction reactions unfavorable; however  the nano scale nature of GNPs could allow for localized reducing envi ronments. GNP showed the ability to survive the oxidation  process. These GNPs would provide the source of carbon needed to reduce Ta2O5 to TaC above 1135 °C; thus suppress expected,  this consumes oxygen, making less oxygen avail ing the formation of the oxide layer. Electron dispersive spec able to penetrate further into TaC. The HR-TEM micrograph  troscopy (EDS)  is performed along the cross section of  the  of the oxide layer in Fig. 11 provides evidence of GNP sealing  oxides in order to examine the amounts of carbon and oxygen  a grain boundary and surviving the plasma flow. The GNP identified by its signature \\x180.35 nm layer spacing,  structure,  in the oxide layers. The EDS results for TaC shown in Fig. 13a,  indicate that throughout the oxide layer the amount of oxy lies on top of a TaO3 grain and an unidentifiable oxide grain  gen exceeds that of carbon. In contrast, EDS results for a typ effectively sealing the grain boundary. The survival of  the  ical TaC-GNP sample shown in Fig. 13b,  indicate that  the  GNP indicates the plausibility of GNP hindering oxygen influx  oxide layer contains more carbon than oxygen. This increase  throughout the oxidation process by both physical and chem in carbon content is believed to be due to the presence of car ical means.  bon from GNPs and that found in TaC. The reduction of oxy The effect of  the grain sealing mechanism is further ex gen is due to the reduction of Ta2O5  to TaC and CO gases  plored by isothermal TGA experiments on TaC and TaC-GNP samples at 1000 °C. The conditions here are not nearly as se which escape the sample. Additional evidence of GNP in duced  localized  reduction  environments  comes  from the  vere and therefore the effects of bursting and increased struc HR-TEM micrograph presented in Fig. 14. Analysis of plane  tural integrity don’t play a role here. Any improvements in the  spacings  reveals  that adjacent carbon and TaC phases are  oxidation behavior would be largely due to GNP grain sealing  embedded  in  the  amorphous  structure.  The  amorphous  mechanisms. The TGA results are shown in Fig. 12 and it can  be seen that all GNP reinforced TaC samples oxidize to a lower  phase is the result of Ta2O5 reaching its melting point of 1872 °C and rapidly solidifying upon cessation of the plasma  extent  than the TaC sample. The conditions here are less  flow exposure. The high temperature needed to melt Ta2O5  Fig. 13 - Thermal gravimetric analysis (TGA) at 1000 °C shows GNP reinforced samples oxidize to a lower extent. GNP grain  sealing is dominant mechanism at this temperature range.  (A color version of this figure can be viewed online.)  \\x0c', '406  C A R B O N 6 7 ( 2 0 1 4 ) 3 9 8 - 4 0 8  the interface of the oxide and the unoxidized sub layer by X ray micro CT. The TaC and TaC-GNP interfaces,  shown in  Fig. 15 as well as in Supplementary video files V1 and V2 (avail able online), can be seen to have vastly different morphologies.  The TaC interface has large protruding interconnected cracks  throughout the surface, while the TaC-GNP sample is largely  crack free. The TaC sample experiences severe cracking mostly  due to thermal shock. These large cracks further enable the  penetration of oxygen thus accelerating the formation of  the  oxide layer and the consumption of the unoxidized material.  The superior structural  integrity of the TaC-GNP interface  region is the result of a combination of the mechanisms dis cussed before. The TaC-GNP samples are inherently tougher  and thus less susceptible to cracking because of GNP tough ening mechanisms such as grain wrapping. The higher ther mal  conductivity  of  TaC-GNP  induces  small  thermal  gradients thereby reducing thermal shock which can initiate  crack nucleation and propagation. Furthermore, the suppres sion of cracking at the interface with the oxide layer inhibits  the penetration of oxygen deeper into the sample. The addi tion of GNP therefore not only suppresses the formation of  the oxide layer,  it protects  the underlying unoxidized sub  layer from the structural damage.  4.  Conclusions  Fig. 14 - TEM micrographs shows presence of TaC near and  graphene structure. Both structures are embedded in  amorphous structures. Amorphous glassy structures are formed when temperatures exceed \\x181100 °C and Ta2O5  melts.  (A color version of this figure can be viewed online.)  also catalyzes the reduction of Ta2O5 upon coming across ex The addition of GNPs is shown to enhance the oxidation resis cess carbon provided by the GNP.  In a molten state, Ta2O5  tance of TaC-GNP composites exposed to a high temperature  would provide a protective layer and thus the newly formed  plasma flow. Through several mechanisms unique to GNP, the  TaC would not readily oxidize. The formed CO gases however  formation of the oxide layer thickness is suppressed by up to  would be expected to easily escape the molten Ta2O5.  3.3.4.  Unoxidized sub layer structural  integrity  60%. The high thermal conductivity of GNPs allows heat  to  flow throughout the sample and away from the front surface  which experiences the most extreme conditions. The tough Aside from the suppression of  the formed oxide layer,  it  is  ening effect of GNPs enables them to resist cracking due to  imperative for a UHTC to maintain its structural integrity dur thermal shock, chemical attack, and the high speed plasma  ing exposure to harsh aerothermal conditions. The oxide layer  flow. The suppression of cracks is vital to hindering the rate  is porous by nature and not expected to provide structural  of oxidation as cracks enable the rapid penetration of oxida strength; however  it  is highly desirable  for  the underlying  tion. SEM and HR-TEM analysis reveals that GNP survive the  unoxidized material  to maintain its structural  integrity. The  extreme oxidizing environment,  thereby signaling that GNP  structure of  the TaC and TaC-GNP samples  is evaluated at  may play an active role throughout the oxidation process.  Fig. 15 - Micro computed tomography (MicroCT) results show the interface between formed oxide layer and unoxidized  sublayer. The structural integrity of (a) TaC is considerably deteriorated compared to (b) TaC-GNP. Large cracks can be seen at  the interface of TaC and the oxide layer in the TaC sample.  \\x0c', 'A GNP grain sealing mechanism is proposed,  in which  GNPs seal grain boundaries and inhibit the influx of oxygen.  Surviving GNP also induce localized reducing environments  by providing excess carbon needed to drive the conversion  of the formed oxide, Ta2O5 to TaC. X-ray MicroCT analysis re veals that GNP oxidation resistance mechanisms also protect  the underlying unoxidized structure from structural damage.  The GNP toughening, grain sealing, and thermal conductivity  mechanisms are inherent to GNPs and could be extended to  other ultrahigh temperature ceramics systems.  Acknowledgements  Authors would like to acknowledge Dr. Ali Sayir, Program  Manager of High Temperature Aerospace Materials at  the  Air Force Office of Scientific Research and FA9550-11-1-0334  and FA9550-12-1-0263 Grants. Dr. Allen Gu and Mr. Bryan Maj krzak at XRADIA are acknowledged for extending the use of  their micro  computed  tomography  (Micro  CT)  facilities.  Authors also acknowledge the support of the Advanced Mate rials  Engineering  Research  Institute  (AMERI)  at  FIU.  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},{
  "_id": 141,
  "PDF": "Oxidation behavior of hot pressed ZrB2-SiC and HfB2-SiC composites.pdf",
  "Text": "['Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Contents lists available at ScienceDirect  Int.  Journal of Refractory Metals and Hard Materials  j o u r n a l h o m e p a g e : w w w. e l s e v i e r . c o m / l o c a t e / I J R M H M  Investigations on synthesis of ZrB2 and development of new composites with HfB2 and TiSi2  J.K. Sonber ⁎, T.S.R. Ch. Murthy, C. Subramanian, Sunil Kumar 1, R.K. Fotedar, A.K. Suri  Materials group, Bhabha Atomic Research Centre, Mumbai,  India-400 085  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 26 February 2010 Accepted 10 June 2010  Keywords:  ZrB2 TiSi2 Composite Synthesis Densification Microstructure Oxidation studies  This paper presents the results of experimental investigations carried out on the synthesis of pure ZrB2 by boron carbide reduction of ZrO2 and densification with the addition of HfB2 and TiSi2. Process parameters and charge composition were optimized to obtain pure ZrB2 powder. Monolithic ZrB2 was hot pressed to full density and characterized. Effects of HfB2 and TiSi2 addition on densification and properties of ZrB2 composites were studied. Four compositions namely monolithic ZrB2, ZrB2 + 10% TiSi2, ZrB2 + 10% TiSi2 + 10% HfB2 and ZrB2 + 10% TiSi2 + 20% HfB2 were prepared by hot pressing. Near theoretical density (99.8%) was obtained in the case of monolithic ZrB2 by hot pressing at 1850 °C and 35 MPa. Addition of 10 wt.% TiSi2 resulted in an equally high density of 98.9% at a lower temperature (1650 °C) and pressure (20 MPa). Similar densities were obtained for ZrB2 + HfB2 mixtures also with TiSi2 under similar conditions. The hardness of monolithic ZrB2 was measured as 23.95 GPa which decreased to 19.45 GPa on addition of 10% TiSi2. With the addition of 10% HfB2 to this composition, the hardness increased to 23.08 GPa, close to that of monolithic ZrB2. Increase of HfB2 content to 20% did not change the hardness value. Fracture toughness of monolithic sample was measured as 3.31 MPa m1/2, which increased to 6.36 MPa m1/2 on addition of 10% TiSi2. With 10% HfB2 addition the value of KIC was measured as 6.44 MPa m1/2, which further improved to 6.59 MPa m1/2 with higher addition of HfB2 (20%). Fracture surface of the dense bodies was examined by scanning electron microscope. Intergranular fracture was found to be a predominant mode in all the samples. Crack propagation in composites has shown considerable deflection indicating high fracture toughness. An oxidation study of ZrB2 composites was carried out at 900 °C in air for 64 h. Specific weight gain vs time plot was obtained and the oxidized surface was examined by XRD and SEM. ZrB2 composites have shown a much better resistance to oxidation as compared to monolithic ZrB2. A protective glassy layer was seen on the oxidized surfaces of the composites.  © 2010 Elsevier Ltd. All rights reserved.  1. Introduction  Zirconium diboride is a leading material in the category of ultra high temperature ceramics (UHTC) due to very high melting point (3245 °C), (57.9 WM−1 K−1), high thermal conductivity good thermal shock resistance, low coefficient of thermal expansion (5.9 × 10−6 °C−1), retention of strength at elevated temperatures and stability in extreme environments [1-3]. ZrB2 is considered a candidate material for hypersonic flight, atmospheric re-entry and rocket propulsion [1,4,5]. It gets wetted but not attacked by molten metals and hence is used for holding molten metal and as thermo-well tubes in metal processing [6]. Good electrical conductivity makes it suitable for electrode application in Hall-Heroult cell and electric discharge machining [7-9]. ZrB2 can be synthesized by (a) reaction between Zr and B [10] (b) borothermic reduction of ZrO2 [11], (c) boron carbide reduction of ZrO2 in the presence of carbon [12], (d) carbothermic reduction of ZrO2 and  ⁎ Corresponding author. Tel.: +91 22 2559 0473; fax: +91 22 2550 5151. E-mail address: jitendra@barc.gov.in (J.K. Sonber). 1 Post Irradiation Examination Division, BARC.  0263-4368/$ - see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2010.06.007  B2O3 [1] (e) metallothermic reduction of ZrO2 and B2O3 [13-15] and (f) chemical vapor deposition. The synthesis of ZrB2 from its elements is uneconomical due to the high cost of both Zr and B powders. Borothermic reduction of ZrO2 to obtain ZrB2 also involves the use of expensive boron powder. In carbothermic reduction of ZrO2 and B2O3, loss of boron occurs due to the evaporation of boron oxides during the reaction and the product is not pure ZrB2 but a mixture of borides. In case of metallothermic reduction of ZrO2 and B2O3 the product gets contaminated with the metal borides of reductant metal. ZrB2 powder can also be synthesized by SHS reaction between ZrO2, B2O3 and Mg. Chemical vapor deposition techniques are suitable for coating and not for bulk production of powders. Preparation of ZrB2 by boron carbide reduction of ZrO2 [12] according to reaction (1) seems to be the best route since it involve use of cheap raw materials and results in pure boride with minimum boron loss.  ZrO2 þ 1=2B4C þ 3=2C→ZrB2 þ 2CO↑  ð1Þ  Despite having excellent properties, the actual uses of ZrB2 are limited due to its poor sinterability and low fracture toughness. Due to  \\x0c', \"22  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  strong covalent bonding and low self-diffusion, high temperature and external pressure are required to densify monolithic ZrB2 [1,10]. Martinez et al. [16] have densified monolithic ZrB2 to 86.5% ρth by hot pressing at 1900 °C and 30 MPa pressure. Chamberlain et al. [17] have obtained a density of  98% ρth by pressureless sintering ZrB2 at 2150 °C for 540 min. The same authors have reported the densification of ZrB2 powder to 99.8% by hot pressing at 1900 °C and 32 MPa [18]. By reactive hot pressing of zirconium and boron powder mixture, 99% dense ZrB2 was obtained at 2100 °C [19]. Several additives have been tried to improve the sinterability and properties of ZrB2. Mishra et al. [20] have reported that oxygen present on the surface of ZrB2 hinders densification. By the addition of TiC and C they have achieved 94% ρth by pressureless sintering at 1800 °C in Ar atmosphere. Zhu et al. [21] have sintered ZrB2 powder to near theoretical density at 1900 °C without the application of external pressure using carbon coated starting powders. Carbon addition was found to effectively remove oxide impurities from the surface of ZrB2 particles. Fahrenholtz et al. [22] have sintered ZrB2 to full density at 1850 °C using a combination of B4C and carbon as additive. Rangaraj et al. [23] have obtained ZrB2-ZrC composite by reactive hot pressing of Zr and boron carbide mixture at 1600 °C and 40 MPa. Apart from these additives, silicon carbide is the most common, as it has been found to enhance the strength and oxidation resistance of zirconium diboride ceramics [1,18]. Yan et al. [24] have densified ZrB2-20 wt.% SiC by pressureless sintering at 2250 °C using 4 wt.% Mo as additive. The composite showed high fracture toughness of 5.39 MPa m1/2 .. Zhu et al. [25] have hot pressed ZrB2-30 vol.% SiC composite to 99.8% ρth at 1900 °C and 32 MPa. The composite exhibits high flexural strength of 900 MPa. M. Zhu et al. [26] have used liquid polycarbosilane (LPCS) as additive to obtain 95% dense ZrB2-SiC composite by pressureless sintering at 1900 °C. Addition of 4% Ni improves the densification and results in 98% density while hot pressing at 1850 °C [16,27] . Monteverde et al. [28] have found that addition of 2.5% Si3N4 causes formation of grain boundary glassy phase and results in 98% ρth density at 1700 °C and 30 MPa. MoSi2 (15 to 20%) was found to help in densification, resulting in 98% dense ZrB2 at 1750 °C and 30 MPa [29,30]. Sciti et al. [31] have densified ZrB2-20 vol.% MoSi2 at 1850 °C and obtained fine grained structure. Recently Shu-Qi Guo et al. [32] have studied the effect of ZrSi2 addition on pressureless sintering of ZrB2 and obtained 99.5% ρth sample at 1600 °C in vacuum by the addition of 20% ZrSi2. As per authors' knowledge there is no literature on the use of TiSi2 and HfB2 as sinter additive to ZrB2. This paper presents the study on  Fig. 2. Thermogravimetry of ZrB2 synthesis.  the effect of TiSi2 and HfB2 addition on processing and properties of ZrB2.  2. Experimental  2.1. Starting material  Raw materials used were ZrO2 (99% purity; 8.34 μm median diameter), boron carbide powder (78.5% B, 19.5% C, b 1% O, 0.02% Fe, 0.02% Si, 5.34 μm median diameter; supplied by M/S Boron 13.9 μm median Carbide India) and petroleum coke (C-99.4%, diameter, supplied by M/S Assam carbon, India). All the raw  Fig. 3. XRD Pattern of thermogravimetry product.  Table 1 Effect of temperature and charge composition.  S. no.  Molar ratio ZrO2:B4C:C  Temperature (°C)  Weight loss (%)  Phases present  Carbon (%)  Oxygen (%)  1. 2. 3. 4. 5. 6. 7.  2:1:3 2:1:3 2:1:3 2:1:3 2:1:3 2:1.1:2.7 2:1.1:2.7  1200 1400 1600 1700 1800 1800 1875  18.47 27.15 30.28 31.66 33.04 32.67 33.01  ZrB2, ZrO2 ZrB2, ZrO2 ZrB2, ZrB, ZrO2 ,C ZrB2, ZrB, C ZrB2, ZrB, C ZrB2 ZrB2  9.7 8.2 7.3 6.1 3.2 1.3 0.06  5.22 2.9 2.4 2 0.57 1.5 0.5  Fig. 1. XRD pattern of starting materials ZrO2, B4C and petroleum coke.  \\x0c\", \"J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  23  2.2. Synthesis  A small pellet of the reactants in stoichiometric quantity (as per reaction 11) was prepared by weighing accurate quantity of individual components, mixing it thoroughly and pelletizing. This pellet was used for thermogravimetric investigations (Setaram TAG 24) up to 1500 °C. For regular synthesis of ZrB2, weighed quantities of zirconium dioxide, boron carbide and petroleum coke in various ratios were mixed thoroughly in motorized mortar and pestle. The powder mixture was then pelletized under pressure of about 280 MPa to obtain pellets of  Fig. 4. Free energy change of reactions (thermodynamic data from Barin's table [34]).  materials were dried in an oven at 100 °C to remove moisture content be fore use . XRD pattern of the raw materials are presented in Fig. 1.  Fig. 5. XRD pattern of the products obtained by varying charge composition. MR: molar ratio (ZrO2:B4C:C).  Fig. 6. Effect of temperature and charge composition on carbon and oxygen content of the product.  Fig. 7. SEM image of powders (a) ZrB2, (b) HfB2 and (c) TiSi2.  \\x0c\", \"24  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Fig. 8. Particle size distribution of powders (a) ZrB2, (b) HfB2, (c) TiSi2 and (d) mixed charge.  20 mm diameter. While mixing and pelletizing, all the materials coming in contact with the powder were made of tungsten carbide. The pellets were then charged in a graphite crucible and heated in an induction furnace under a dynamic vacuum of 2 × 10−5 mbar and at fixed temperature between 1200 and 1875 °C for 2 h. Temperature of the charge was measured using a two-colour pyrometer with an accuracy of ± 22 °C. After completion of the reaction, the furnace was cooled to room temperature in vacuum and the reacted pellets were taken out, crushed and ground to fine size using high-energy cup grinding mill with tungsten carbide lining. The major phases of the powders were identified by XRD (Cu Kα (λ = 1.5404 Å radiation in a Philips PW1830 diffractometer) and impurities were analyzed by chemical methods. The median particle diameter and particle size distribution were measured using laser particle size analyzer (CILAS PSA 1064 L). SEM (20 kV, Philips, FEI XL30) of the powders was also used to cross check the size and morphology.  2.3. Densification and characterization  For densification, weighed quantities of fine zirconium diboride, hafnium diboride and titanium disilicide were mixed thoroughly using a motorized mortar and pestle in dry condition for 1 h and powder mixture of four different compositions were prepared as (1) ZrB2 (2) ZrB2 + 10% TiSi2, (3) ZrB2 + 10% TiSi2 + 10% HfB2 (4) ZrB2 + 10% TiSi2 +  Table 2 Effect of sinter additives on densification and properties of ZrB2.  Additive  Temp. (°C)  Pressure (MPa)  Density (%)  Hardness (GPa)  KIC  (MPa m1/2)  nil 10 wt.% TiSi2 10% HfB2 + 10 wt.% TiSi2 20% HfB2 + 10 wt.% TiSi2  1850 1650 1650 1650  35 20 20 20  99.8 98.9 99.6 98.4  23.91 ± 1.5 19.45 ± 1.9 23.08 ± 1.3 23.66 ± 1.8  3.31 ± 0.2 6.36 ± 1.0 6.44 ± 1.1 6.59 ± 0.7  20% HfB2. The powders were then loaded in a high density graphite die (12 mm hole) and hot pressed at temperatures of 1650 °C to 1850 °C under a pressure of 20 to 35 MPa for 60 min in a high vacuum (1 × 10−5 mbar) chamber. The pellets were ejected from the die after cooling and the density measured by Archimedes' principle. Densified samples were polished to mirror finish using diamond powder of various grades from 15 to 0.25 μm in an auto polisher (laboforce-3, Struers). Microhardness was measured on the polished surface at a load of 100 g and dwell time of 10 s. The indentation fracture toughness (KIC) data were evaluated by crack length measurement of the crack pattern formed around Vickers indents (using 10 kg load), adopting the model formulation proposed by Anstis et al. [33] KIC = 0.016(E/H)1/2P/c3/2, where E is the Young's modulus, H the Vickers hardness, P the applied indentation load, and c  Fig. 9. XRD Pattern of ZrB2 based composites.  \\x0c\", 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  25  Table 3 Oxidation parameters for ZrB2 composites.  Sample  Monolithic ZrB2 ZrB2 + 10 wt.% TiSi2 ZrB2 + 10% HfB2 + 10 wt.% TiSi2 ZrB2 + 20% HfB2 + 10 wt.% TiSi2  m  0.97 2.52 2.38 2.02  Km  4.5 × 10−7 1.96 × 10−10 1.67 × 10−10 2.03 × 10−10  Kp (Kg m−4 s−1) × 108  1.25 × 10−8 1.80 × 10−9 9.53 × 10−10 2.22 × 10−10  the half crack length. The reported values of hardness and fracture toughness are the average of five measured values. Fractured surface of dense pel lets was analyzed by scanning electron microscope.  2.4. Oxidation  Hot pressed pellet of diameter 12 mm was cut into thin slice of 3 mm thickness by high speed diamond cutter. All the surfaces of the cut sample were polished with emery papers (1/0, 2/0, 3/0, 4/0) and finally with diamond paste up to 1 μm finish. Oxidation tests were conducted in a resistance heated furnace. In order to avoid oxidation during heating, the sample was directly inserted into the furnace after the furnace temperature reached 900 °C. Samples were placed in an alumina crucible kept into the furnace. The sample was oxidized for different time intervals (0.5, 1, 2, 4, 8, 16, 32, and 64 h) at 900 °C. The sample was carefully weighed before and after exposure, to determine the weight change during the oxidation process. The oxidation products were identified using XRD. The morphology and nature of oxide layer were understood by observing the surface in a scanning electron microscope (SEM).  3. Results and discussion  3.1. Synthesis  Fig. 10. SEM image of ZrB2 + 10%TiSi2 + 20%HfB2 composite and elemental analysis of different phases.  Fig. 2 presents the weight loss vs. time curve obtained by thermogravimetric experiment of stoichiometric charge mixture as per reaction 1. The reaction starts at 1200 °C and results in a total  Fig. 11. Elemental mapping of Zr, Ti and Si  in ZrB2 + 10%TiSi2 + 20%HfB2 composite.  \\x0c', '26  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  loss of  27.5%. The weight weight loss observed is lower than the theoretical weight loss of 33.0%. The XRD pattern revealed the presence of ZrB2, ZrB and graphite phases (Fig. 3). In the case of synthesis experiments carried out in an induction furnace at 1200 °C, the weight loss was 18.47% and the product was impure, containing ZrB2 and ZrO2 phases (Table 1). The carbon and oxygen contents were 9.7 and 5.22% respectively. At 1400 °C, the weight loss increased to 27.15% and carbon and oxygen contents were less by 8.2 and 2.9%. At 1600 °C, weight loss increases to 30.28% but the product contained ZrB2, ZrB, ZrO2 and C. At temperatures of 1700 °C and above, ZrO2 was absent in the product but ZrB was seen. At 1800 °C the weight loss was close to a theoretical value of 33.0%, however the product was  Fig. 12. Fracture surfaces of (a) monolithic ZrB2 (b) ZrB2 + 10%TiSi2 (c) ZrB2 + 10%TiSi2 + 10%HfB2.  composed of ZrB2, ZrB, and carbon. The presence of boron deficient phase (ZrB) even after the treatment at 1800 °C indicates the loss of boron from the charge, which could occur by following reaction.  ZrO2 þ 3=4B4C ¼ ZrB2 þ 1=2B2O3 þ 1=2CO þ 1=4C  ð2Þ  The thermodynamic calculation [34] (Fig. 4) indicates that both reactions (1) and (2) could take place on heating and hence the loss of boron by evaporation as B2O3 during the reaction. Hence to obtain a single-phase ZrB2 phase, it is necessary to add an excess boron in the charge. Fig. 5 presents the XRD patterns of product obtained by stoichiometric and modified charge mixture. It evidently tells that ZrB and graphite phase are not present and single phase ZrB2 is obtained by reacting the modified charge mixture (molar ratio, ZrO2:B4C:C, 2:1.1:2.7) at 1800 °C. The oxygen and carbon contents were found to be 1.5 and 1.3% respectively. This product was further purified to low levels of oxygen (0.06%) and carbon (0.5%) by heating to 1875 °C and soaking at this temperature in vacuum for 30 min. The progressive removal of carbon and oxygen from the product with reaction temperature and the final purification using modified charge are presented in Fig. 6. Hong Zhao et al. [12] have also reported the intermediate reaction resulting in the loss of boron as B2O3.  3.2. Densification and characterization  SEM images (Fig. 7) of the powder shows that ZrB2 and HfB2 particles are of  2-3 μm whereas TiSi2 particles are large at 15-20 μm size. Fig. 8 presents the particle size distribution of ZrB2, HfB2, TiSi2 and mixed charge. It shows that particle size distribution is monomodal for ZrB2 and HfB2 whereas bimodal for TiSi2. During mixing, size reduction of the particles also takes place and the particle size distribution of the mixed powders is narrow and monomodal. Median diameter of the mixed charge was measured as 2.67 μm. Table 2 presents the effects of hot pressing parameters on the density and mechanical properties of the pellets. In the case of monolithic ZrB2, a near full density (99.8% TD) was obtained at 1850 °C with a pressure of 35 MPa and dwell time of 60 min. Addition of 10 wt.% TiSi2 resulted in densification of 98.9% TD at a relatively low tempera ture o f 1650 °C and a low pressure o f 20 MPa . Composites with 10 and 20% HfB2 content were also hot pressed to nearly full density at 1650 °C and 20 MPa . The enhanced sintering at lower hot pressing temperature and pressure is due to the liquid phase sintering with low melting (1540 °C) additive TiSi2 and the low melting reaction product ZrSi2 (1620 °C). XRD pattern of the dense pellets are shown in Fig. 9. All the three samples indicate the presence of crystalline ZrB2 and ZrSi2. ZrSi2 is formed during sintering by the following reaction.  ZrB2 þ TiSi2→ZrSi2 þ TiB2  ð3Þ  TiB2 is not seen in the XRD pattern of the sintered product, probably due to the formation of ZrB2-TiB2 solid solution. HfB2 also forms a solid solution with ZrB2 and hence not seen as a distinct phase in XRD pattern. Post et al. [35] have reported that, these borides have complete mutual solubility. Compar ison o f o ther s in ter add i t ives f rom l i tera ture is discussed below. The addition of 20% MoSi2 results in 98.1% TD on hot pressing at 1800 °C and 30 MPa [30]. 20% SiC addition resulted in near theoretical density on hot pressing at 2000 °C and 30 MPa [36]. Guo et al. [37] have reported that addition of 5% Re2O3 (Re = Y,Yb,La,Nd) to ZrB2-20% SiC results in N 99% TD on hot pressing at 1900 °C. Wang et al. [38] have reported that 10% Mo addition gives a density of 98.9% TD on hot pressing at 1950 °C and 20 MPa . Zhu et al. [39] have reported that addition of 3-10%  \\x0c', 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  27  Fig. 13. Crack propagation in (a) monolithic ZrB2 (b) ZrB2 + 10%TiSi2 (c) ZrB2 + 10%TiSi2 + 10%HfB2 (d) ZrB2 + 10%TiSi2 + 20%HfB2.  Al2O3 and Y2O3 and 20% SiCw to ZrB2 gives a density N 97% TD on hot pressing at 1800 °C. From the above it is clear that the addition of TiSi2 in the present investigation is found very effective in lowering the sintering temperature to 1650 °C. The addition of TiSi2 to TiB2 and HfB2 has also shown similar results [40,41].  3.3. Mechanical properties  Variations in Vickers hardness and fracture toughness of ZrB2 composites are presented in Table 3. Hardness of monolithic sample was measured as 23.91 GPa, which reduced to 19.45 GPa with the addition of 10 wt.% TiSi2. The decrease in hardness is due to the lower hardness of TiSi2 (8-10 GPa) [40]. The effect of TiSi2 is nullified with HfB2 addition. The hardness value composite with 10% HfB2 addition was 23.08 GPa, which marginally increased further to 23.66 GPa with 20 wt.% HfB2. The increase in hardness could be attributed to (Zr-Hf)B2 solid solution formation. Hardness of monolithic ZrB2 has been reported to be in the range of 22-23 GPa [1,2,18,42]. Akin et al. [43] have reported the hardness of ZrB2-20-60% SiC composite to be 26 GPa. Fracture toughness of monolithic ZrB2 sample was measured as 3.31 MPa m1/2, which increased considerably to 6.36 MPa m1/2 with 10% TiSi2 addition. With further addition of HfB2, the fracture toughness of ZrB2 + T iS i2 + H fB2 compos ite was found to be marginally high at 6.44-6.59 MPa m1/2. Fracture toughness of monolithic ZrB2 has been reported as 3.5 MPa m1/2[18]. Enhancement of fracture toughness from 4.52 to 7.98 MPa m1/2 by 5% Mo addition is reported by Wang et al. [38].Zhu et al. [39] have reported the fracture toughness of 6.7 MPa m1/2 for ZrB2 + 20% SiCw + 3%  YAG. Similar value of fracture toughness was reported for ZrB2-SiC nanocomposite prepared by using SiC nanopowder [44].  3.4. Microstructure evolution  Fig. 10 presents the secondary electron image of ZrB2 + 10% TiSi2 + 20% HfB2 composites. It shows the presence of dark phase in the gray matrix. EDS pattern of the phases are also inserted into the picture.  Fig. 14. Specific weight gain vs time plot for ZrB2 based composites. (Temp: 900 °C).  \\x0c', '28  J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  Gray matrix is essentially ZrB2 in which Ti and Hf have diffused whereas dark phase contains mainly Si with some Zr. Fig. 11 presents the elemental mapping for Zr, Ti and Si. It shows that Zr and Ti are distributed uniformly whereas Si is present only in the black phase. Fig. 12 (a-c) presents the fracture surfaces of monolithic ZrB2, ZrB2 + 10% TiSi2 and ZrB2 + 10% TiSi2 + 10% HfB2 composites. The mode of fracture is seen to be intergranular in all the samples. Regular faceted grains are visible. Fig. 13 (a) presents the features of indentation crack in monolithic ZrB2 and Fig. 13 (b-d) in composites. The crack propagation line is almost straight in monolithic ZrB2 whereas considerable deflections are observed in the composite samples, which explain the high fracture toughness.  3.5. Oxidation study  The weight gain data obtained during oxidation at 900 °C as a function of time for ZrB2 composites are presented in Fig. 14. Continuous weight gain with time is observed in all the samples. In the case of composites, the rate of oxidation was found to decrease with increase in time which indicates the formation of protective layer. In monolithic ZrB2, oxidation rate was found to be constant. In order to understand the nature of oxidation, the oxidation data was fitted in the general rate equation for all the composites.  ð  Þm = Km :t Δw = A  ð4Þ  where Δw is the change in weight, A—surface area of the sample, t— oxidation time and Km—rate constant. Km and m values are presented in Table 3. The composite samples have shown a parabolic pattern of oxidation whereas monolithic has shown linear oxidation. Better oxidation resistance of the composite samples is attributed to the formation of silica based glassy layer. The SEM microstructures (Fig. 15) of oxidized surfaces evidently show the formation of protective glassy phase. The glassy phase was analyzed by EDS to contain mainly silicon ( 46 at.%) and oxygen ( 52 at.%). Zirconium ( 0.8 at.%) and titanium ( 0.2 at.%) are also present in very small quantity. A typical EDS pattern from the oxidized surface is shown in Fig. 16. Formation of protective borosilicate glass has been reported in literature [45,46]. The glass formed in this study could be borosilicate glass. EDS analysis has some limitations for detection of boron. Fig. 17 presents the XRD pattern of the oxidized surface. The major crystalline phase in all the composites is confirmed as ZrO2. Peaks of ZrSiO4 and TiO2 are also present. The following reactions are possible during the oxidation process.  2=5ZrB2 þ O2→2=5ZrO2 þ 2=5B2O3  2=5TiB2 þ O2→2=5TiO2 þ 2=5B2O3 1=3ZrSi2 þ O2→1=3ZrO2 þ 2=3SiO2  ZrO2 þ SiO2→ZrSiO4  ð5Þ  ð6Þ  ð7Þ  ð8Þ  Opeka et al. [47] have reported that SiC-containing ZrB2 ceramics had high oxidation resistance up to 1500 °C compared to pure ZrB2 ceramics. Monteverde et al. [27] have reported that of ZrB2 + 5% Ni composites get degraded on oxidation at 1000 °C due to fast oxidation at the Ni rich grain boundaries. Monteverde et al. [48] have reported that oxidation of monolithic ZrB2 starts at 420 °C. It is observed that the introduction of SiC particles markedly improves oxidation resistance due to the formation of an adherent and protective borosilicate glass layer that coats the sample surface, effectively limiting the inward diffusion of oxygen toward the reaction interface. Guo et al. [49] have reported that oxidation of ZrB2 powder follows para-linear kinetics in air at 650-800 °C, where the dominating term is the parabolic one, accounting for oxygen diffusion in the oxide scale.  Fig. 15. SEM images of the oxidized (900 °C, 64 h) surface of ZrB2 with (a) 10%TiSi2, (b) 10%TiSi2 + 10%HfB2, and (c) 10%TiSi2 + 20%HfB2.  Guo et al. [50] have observed that addition of Si improves the oxidation resistance of ZrB2 at 1500 °C whereas Zr addition decreases the oxidation resistance. Karlsdottir et al. [46] have reported the formation of ZrO2 islands in the borosilicate glass on oxidation of ZrB2 + SiC ceramic at 1500 °C.  4. Conclusion  Charge composition and synthesis parameters were optimized to obtain pure ZrB2 powder by boron carbide reduction of ZrO2. Monolithic  \\x0c', 'J.K. Sonber et al.  /  Int.  Journal of Refractory Metals and Hard Materials 29 (2011) 21-30  29  Fig. 16. Typical EDS pattern of the oxidized surface in ZrB2 + 10% TiSi2 + 20% HfB2 composites.  [5]  Fig. 17. XRD pattern of the surface oxidized at 900 °C for 64 h in air.  ZrB2 was hot pressed to full density at 1850 °C and 35 MPa. Addition of 10 wt.% TiSi2 to ZrB2 resulted in an equally high density of 98.9% by hot pressing at a lower temperature (1650 °C) and pressure (20 MPa). Similar densities were obtained for ZrB2 + HfB2 mixtures also. Hardness of monolithic sample is 23.95 GPa which decreases to 19.45 GPa on addition of 10% TiSi2. On further addition of 10% HfB2, the hardness increased to 23.08 GPa, which is very close to that of monolithic ZrB2. Fracture toughness of monolithic sample is 3.31 MPa m1/2, which increases to 6.36 MPa m1/2 on addition of 10% TiSi2. Further addition of HfB2 marginally increased the value of KIC to 6.44 and 6.59 MPa m1/2 with 10% and 20% HfB2 addition respectively. Fracture mode was found to be intergranular. Crack propagation in composites has shown considerable deflection indicating high fracture toughness. TiSi2 has been found to be a useful additive in sintering ZrB2. This helps in reducing the hot pressing temperature by 200 °C. This composite and also the addition of HfB2 have been found to have a better oxidation resistance at 900 °C. The glassy phase on the oxidized surface has been found to be a mixture of silica and silicates.  References  [1]  Fahrenholtz WG, Hilmas GE. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90(5):1347-64. [2] Bauccio ML. ASM engineered materials reference book. United States of America: ASM International; 1994. 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},{
  "_id": 142,
  "PDF": "Oxidation Behavior of SiC Platelet-Reinforced ZrB2 Ceramic Matrix Composites- Oxidation Behavior of ZrB2 Ceramic Matrix Composites.pdf",
  "Text": "['Int. J. Appl. Ceram. Technol., 9 [1] 178-185 (2012)  DOI:10.1111/j.1744-7402.2010.02647.x  Ceramic Product Development and Commercialization  Oxidation Behavior of SiC Platelet-Reinforced ZrB2 Ceramic Matrix Composites  Mingfu Wang*  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, People’s Republic of China  State Key Lab of New Ceramics and Fine Processing, Department of Materials Science and Engineering,  Tsinghua University, Beijing 100084, People’s Republic of China  Chang-An Wang  State Key Lab of New Ceramics and Fine Processing, Department of Materials Science and Engineering,  Tsinghua University, Beijing 100084, People’s Republic of China  Lei Yu  State Key Lab of New Ceramics and Fine Processing, Department of Materials Science and Engineering,  Tsinghua University, Beijing 100084, People’s Republic of China  School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing  100083, People’s Republic of China  Yong Huang  State Key Lab of New Ceramics and Fine Processing, Department of Materials Science and Engineering,  Tsinghua University, Beijing 100084, People’s Republic of China  Xinghong Zhang  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, People’s Republic of China  This work was  supported by the 973 program from Ministry of Science and Technology of China (Grant No: 5133102-4).  *wmfhit@sina.com.  r 2011 The American Ceramic Society  \\x0c', 'www.ceramics.org/ACT  Oxidation Behavior of ZrB2 Ceramic Matrix Composites  179  Zirconium diboride (ZrB2) ceramic matrix composites reinforced by silicon carbide platelets (SiCpl) were prepared by hot pressing. Five groups of specimens were prepared, in which the contents of SiC platelets were 0, 5, 10, 15 and 20 vol%, respectively. The oxidation behaviors of specimens at different temperatures from 8001 to 16001C for 1 h and at 12001C for  times  different  from 0.5 to 4 h were investigated. With the increase of SiCpl content, trend. Oxidation thermodynamics and kinetics of the SiCpl/ZrB2 composites were discussed. A characteristic hollow quadrangular prismy ZrO2 crystal emerged on the sample surface after oxidation was found and the possible formation mechanism was also presented. The surface element analysis indicated that the main products of oxidation were m-ZrO2, SiO2, and borosilicate (aluminosilicate).  the mass gain exhibited a decreasing  structure of  Introduction  As  the development of aeronautic and astronautic  industry advances, much higher requirements have been  laid on structural materials used in the  extreme  envi ronment for applications such as hypersonic space shut platelets  into ZrB2 chanical properties and oxidation resistance of the ZrB2 ceramics.  ceramics would improve  the me In the present paper, SiC platelets were added into  the ZrB2 matrix with AlN as sintering aid. SiCpl/ZrB2 ceramic matrix composites were subsequently prepared  tle,  the  edge of  reentry  guided missile,  and so forth.  by hot pressing. The mechanical properties and oxida Much attention has been focused on ultra high temper tion behavior of  the  specimens were  investigated, and  ature ceramic materials (UHTCs) in the recent decades.  the oxidation mechanisms were also discussed.  Zirconium diboride (ZrB2) is an ideal candidate for the UHTCs because of its high melting temperature, low  density, high hardness, high thermal conductivity, and  medium coefficient of thermal expansion. However, the  brittleness  at  low temperature  and weak oxidation re sistance at high temperature limit applications of ZrB2 ceramics. Therefore, much more research has been car ried out to reinforce ZrB2 ceramics and improve oxidation and ablation resistance. Adding a secondary phase  has been proved to be effective in improving the oxida tion resistance as well as the mechanical Kaufman et al.1 firstly demonstrated that  properties.  the addition  of silicon carbide (SiC) improved the mechanical prop erties  and  oxidation  resistance  of metal  diboride  in  1970s; similar conclusions have been obtained by MonBellosi.2 colleagues3-5  Fahrenholtz  teverde  and  and  Experimental Procedure  Materials  Commercially available ZrB2 powder (Gongyi Ceramic Material, Henan, China) with d50 5 2.3 mm and purity \\x15 98% and SiC platelets (Dongfang Ceramic Material, Ningxia, China) with thickness 3-4 mm and length 20 mm (Fig. 1) were used. The major impurity of ZrB2 was ZrO2 and the amount was o2%. AlN  studied  that  a  series  of ZrB2-SiC composites the properties of the composites were  and  found  improved  obviously. SiC, ZrO2, Si3N4, and AlN were added into ZrB2and HfB2-based ceramics in previous research6-8 with excellent results achieved. The oxidation mecha nisms of different composites were also analyzed. The SiC reinforcements such as whiskers,9 particles,10 fibers11 were used in previous  studies, but using SiC  or  platelets (SiCpl) as reinforcements for ZrB2 ceramics was rarely reported. SiC platelets were successfully used as  reinforcements for some ceramics and Si3N4,14 and obvious tained. Therefore, it is expected that  including Al2O3 reinforcing effects were ob 12,13  incorporating SiC  Fig. 1.  SEM photo of SiC platelets.  \\x0c', '180  International Journal of Applied Ceramic Technology—Wang, et al.  Vol. 9, No. 1, 2012  was  added as  sintering aid with 5 vol% (volume  frac used to analyze the phase change of the oxidized and as tion). The SiCpl additions in the ZrB2 matrix were 0, 5, 10, 15, and 20 vol%, respectively. The specimens were  sintered specimens. The surface and cross-section of the  specimens after oxidation were analyzed by using a scan denoted as ZB, ZS1, ZS2, ZS3, and ZS4, respectively.  ning  electron microscope  (SEM, Shimadzu SSX-550)  The powders were mixed by ball milling using zirconia  and energy-dispersive microanalysis  (EDX, Shimadzu,  balls in absolute ethanol for 4 h on planetary mill. Most  Tokyo, Japan).  of SiC platelets maintain the original morphology after  ball milling. The mixture was then dried and sieved with  a 100-mesh screen. Hot-pressing sintering was carried out at 19501C for 1 h in flowing argon atmosphere with  a pressure of 30 MPa.  Results and Discussion  Mechanical Properties of As-Sintered  SiCpl/ZrB2 Composites  Characterization and Analysis  Density of  the sintered ceramics was measured by  The mechanical properties of the SiCpl/ZrB2 composites are shown in Table I.15 It could be seen that the  Archimedes method, while  the  theoretical density was  addition of SiC platelets improved the relative density of  estimated with the  rule of mixture. Bending strength  was determined by three-point bend testing (test bars 4 mm \\x02 3 mm \\x02 36 mm) with a span of 30 mm and a crosshead speed of 0.5 mm/min. Fracture toughness was  measured by single-edged notch beam method bars 4 mm \\x02 6 mm \\x02 30 mm) with a and a crosshead speed of 0.05 mm/min, the width of the notch being o0.25 mm. Five specimens were tested each condition. Specimens of 3 mm \\x02 4 mm \\x02 under 20 mm were prepared for oxidation test. All specimens  span of 24 mm  (test  the SiCpl/ZrB2 ceramic matrix composites. It suggested that the addition of SiCpl was beneficial for the densification of ZrB2 since SiC removed the surface oxidation layer of ZrB2 particles and improved the sinterability of ZrB2 materials.16 With the SiC platelet content, fracture toughness of  increase of  improved  the SiCpl/ remarkably and  composites  ZrB2 reached a maximum value  were  even  though  the  bending  at 15 vol% SiCpl strength showed  addition  a weak  decline  trend. More  energy  released when the  cracks  were cleaned in ethanol bath using ultrasonic, dried and  pass through the platelets and the toughness improved.  weighted. In order to study the effects of temperature on  oxidation behavior of ZrB2-based composites, the specimens were tested at different temperatures from 8001 to 16001C with holding time 1 h (after the temperature  achieved the test temperature). For kinetics study, the oxidation experiments were carried out at 12001C with  However, excessive addition of SiCpl inhibited the densification of ZrB2 materials and the flexural strength and hardness decreased because the dimension of SiC plate lets was  somewhat  larger  and induced flaws  into the  ZrB2 matrix. Figure 2 shows  the polished surface morphology  different holding time from 0.5 to 4 h. Three parallel  before oxidation. The dark area as  the arrow directed  specimens were tested for each oxidation condition. The  specimens were heated in an air-atmosphere furnace on an Al2O3 plate with a heating rate of 51C/min. X-ray diffraction (Shimadzu XRD-7000, Tokyo, Japan) was  was SiC platelet, and the undertint area was ZrB2 phase. As can be seen from Fig. 2, the ZrB2 grains were tightly interconnected with uniform grain size and the SiCpl was homogeneously distributed in the matrix.  Samp le  RD (% )  Flexura l strength (MPa)  Fracture toughness (MPa m1 /2)  Hardness (GPa )  Table I.  Mechanical Properties of  the SiCpl/ZrB2 Composites  ZB  ZS1  ZS2  ZS3  ZS4  94.0  97.4  98.0  99.0  97.1  576.84725.83 625.34721.46 517.43725.14 522.25749.46 443.27711.31  6.4070.31 6.8370.42 7.1070.43 8.3570.26 7.4070.20  10.770.61 14.6070.84 13.7970.82 13.8971.23 11.2870.57  \\x0c', 'www.ceramics.org/ACT  Oxidation Behavior of ZrB2 Ceramic Matrix Composites  layer  181  formed  below  exceeded  and the  the  research also showed that  compact SiO2 base materials from oxidation  to protect 16001C, however, when the temperature 16001C the oxidation layer became complex protection layer of SiO2 disappeared.17,18  Fig. 2.  SEM photo of  the SiCpl/ZrB2 composites (arrows point  to  SiC platelets).  Effect of SiCp l Content and Oxidation Temperature on Oxidation Behavior of SiCp l/ZrB2 Composites  The oxidation behavior of SiCpl/ZrB2 composites with different SiCpl contents is shown in Fig. 3. It could be seen that the addition of SiC platelets was beneficial  for the oxidation resistance of composites distinctly and  the mass gain decreased with increasing SiCpl content. The oxidation resistance of SiCpl/ZrB2 composites was improved due to the formation of silica layer, which  prevented  oxygen entering  the  interface  between the  underlying  material  and  oxidation  layer.  Similar  Effect of Oxidation Time on the Oxidation Behavior  of SiCpl/ZrB2 Composites  The specimen ZS3 with good mechanical proper ties was  selected for  studying the oxidation kinetics of  the SiCpl/ZrB2 composites. The mass change of sample ZS3 during oxidation at 12001C with different time is  shown in Fig. 4. The mass change was generally linear at  the beginning of oxidation and then exhibited a para bolic curve relation with the variation of time. Different  rules of mass change were attributed to different oxida tion mechanisms. At  the  incipient  stage of oxidation,  contact oxidation dominated the  reaction because  the  surface material of  the  specimen reacted with oxygen  directly, so the curve was linear. After an oxidation layer  formed  at  the  surface,  the  oxygen  reacted with  the  underlying material after diffusion through the oxidation  layer. In this  stage,  the diffusion of oxygen, which obeys  the Second Fick’s Law, controlled the oxidation behavior,  so the curve matched well with the parabolic type.  Phase Analysis and XRD Patterns after Oxidation  The XRD patterns of  the SiCpl/ZrB2 with different SiCpl content after oxidation are shown in  composites  Fig. 3.  The oxidation behavior of SiCpl/ZrB2 composites with different SiCpl contents.  Fig. 4. The oxidation behavior of the SiCpl/ZrB2 composites with different oxidation time at 12001C in the air.  \\x0c', '182  International Journal of Applied Ceramic Technology—Wang, et al.  Vol. 9, No. 1, 2012  Fig. 5.  X-ray diffraction patterns of  the SiCpl/ZrB2 specimen  before and after oxidation.  Fig. 6.  DG-T curves of  the possible reactions during the  oxidation.  Fig. 5. The XRD pattern of ZS2 before oxidation was  idation product  also shown in Fig. 5. It was observed that the main oxat 12001C was m-ZrO2. Aluminum silicate was also generated during the oxidation. As the  SiCpl content  the SiO2 content accrued obviously by the analysis of EDX. When there was no SiC  increased,  platelet addition,  in specimen ZS0, Al2B4O36 was generated by the reaction of B2O3 and Al2O3. Al2O3 was from the oxidation of sintering aid AlN. Through the  calculation of Gibbs-free energy of reactions as shown in  the following reactions,  it could be seen that AlN was  easily oxidized. More SiO2 was SiC platelet reacting with Al2O3, and then the reaction resulted in Al2SiO5. It was also noted that more Al2SiO5 generated with the increase of SiC platelet.  generated with more  There probably  are  a  few oxidation reactions  as  follows:  ZrB2 þ 5=2O2 ! ZrO2 þ B2O3  2AlN þ B2O3 ! 2BN þ Al2O3  SiC þ 2O2 ! SiO2 þ CO2  SiC þ 3=2 O2 ! SiO2 þ CO  SiO2 þ ZrO2 ! ZrSiO4  4AlN þ 7O2 ! 4NO2 þ 2Al2O3  SiO2 þ Al2O3 ! Al2SiO5  ð1Þ  ð2Þ  ð3Þ  ð4Þ  ð5Þ  ð6Þ  ð7Þ  The  changes  of  the Gibbs-free  energy  of  all  the  above reactions at different temperatures were calculated  as  shown in Fig. 6 (all of  the original data were  after  NIST-JANAF Thermochemical Tables). Reaction (1)  occurred easily because of  the  lowest  reaction Gibbs free  energy. The major oxidation products were ZrO2 and B2O3, and the oxidation of SiC also began at about 12001C. It was consistent with the results of phase anal ysis by XRD. Owing to the rapid volatilization of B2O3 over 11001C,19 the main phase of the reaction product was m-ZrO2. With increasing SiCpl content, SiO2 from the product of reaction (3) covered the material surface  homogeneously and formed an effective oxygen isolating  layer to prevent  the material  from further oxidation.  Analysis of Specimen Surface after Oxidation  Figure  7  shows  the  surface morphology  of  the  SiCpl/ZrB2 composites with different SiCpl after oxidized at 12001C. Two phenomena should be  contents  noted:  (1) many hollow quadrangular prismy  crystals  emerged on the  SiCpl/ZrB2 composites and the quadrangular prismy crystals tended  oxidation surface  of  the  to be perpendicular  to the  surface of  the  composites;  (2) with increasing SiCpl content, the size and content of the quadrangular prismy crystals decreased significantly.  According to XRD and EDX analysis, the quadrangular  prismy crystals were recognized as m-ZrO2. In order to explain the growth behavior of the qua drangular prismy m-ZrO2 was presented as shown in Fig. 8. Exposed in the air at  crystals, a schematic model  \\x0c', 'www.ceramics.org/ACT  Oxidation Behavior of ZrB2 Ceramic Matrix Composites  183  Fig. 7.  Surface morphology of  the SiCpl/ZrB2 composites after oxidation at 12001C for 1 h.  12001C,  the ZrB2-based ceramics were oxidized, the oxidation products such as ZrO2, SiO2, and B2O3 surface of the samples. ZrO2  generated on the  were  and  easily  formed  crystalline  nuclei  on  the  surface  and  growth toward the environment  through mass  transfer  and diffusion. B2O3 easily volatilized because of its high vapor pressure above 11001C. Firstly, several small  B2O3 With  the  bubbles  nearby  gathered  and  then  volatilized.  volatilization of B2O3, B2O3 wrapped with ZrO2, provided a driving force for ZrO2 crystal growth along the easily growing direction. The  bubbles,  en quadrangular  prismy m-ZrO2 Meanwhile, B2O3 went on volatilizing along the ZrO2 tube. With the mass transfer and diffusion carrying on,  formed.  crystals  tubes  Fig. 8.  Schematic diagram of  formation process of  the hollow  quadrangular prismy ZrO2 crystals.  more ZrO2 moved to the bottom of the ZrO2 tube grew longer. The ZrO2 tubes were not homogeneously distributed on the whole surface but  crystal and then  only at some local places, which confirmed the gather of  B2O3 bubbles before the ZrO2 tubes  formed.  \\x0c', '184  International Journal of Applied Ceramic Technology—Wang, et al.  Vol. 9, No. 1, 2012  In order  to figure out  the orientation of  the qua drangular prismy m-ZrO2 the oxidation surface were determined with an XRD.  the pole figures of  crystals,  Figure 9(a) shows that the strongest diffraction peak of the test sample appeared at the angle of 34.21 for (002) crystal plane of m-ZrO2, while the strongest diffraction peak in the standard powder diffraction card of m-ZrO2 \\x16111. is (002) diffraction peak is the third strongest  diffraction peak. Therefore, a preferred growth orienta tion  existed  in  the  formation  product  on the  surface  of  the  of m-ZrO2 samples. Figures  oxidation  9(b)  and (c) were pole figures of the oxidation surface along ½\\x16111\\x8aand [002] directions, respectively. The partly grain growth originally along ½\\x16111\\x8a deflected to the [002] orientation during the oxidation. The results showed that  partial m-ZrO2 [002].  grain grew in a preferred orientation  With increasing SiCpl content, more SiO2 was produced and an amorphous glass layer mainly containing  borosilicate  finally  formed,  preventing  oxygen  from  further  reacting with underlying material and reduced  the speed of oxidation.  Figure 10 presents the cross-section of the specimen after oxidized at 12001C for 1 h and the dashed line is  Fig. 9. The XRD patterns of ZS0 after oxidation (a) and the pole figure along the crystal orientation ½\\x1611 1\\x8a  (b) and [0 0 2]  (c).  Fig. 10.  Cross-section of  the SiCpl/ZrB2composite specimen after oxidized at 12001C for 1 h.  \\x0c', 'the oxidation layer  thickness marker. The thickness of about 50 mm when the SiCpl content was 5 vol% and the thickness of the oxidation  the oxidation layer was  layer was much less  than that of  the pure ZrB2 the same temperature. As the SiCpl content increased, the thickness of oxidation layer decreased and reached as thin as 10 mm when the SiCpl was 20 vol%. It was thus obtained that the borosilicate (aluminosilicate)  after  oxidized at  formed in the material  surface slowed down the diffu sion of oxygen effectively and consequently reduced the  thickness of oxidation layer significantly. The similar also reported elsewhere.20 The oxidation  results were  resistance  of  SiCpl/ZrB2 addition of SiC platelets.  improved  greatly with  the  Conclusions  The addition of SiC platelets improved the relative  density  and mechanical properties  of  the  SiCpl/ZrB2 composites with the maximums reached when the SiCpl content was 15 vol%. The oxidation test at 12001C for  1 h of  the composites showed that  the mass change de creased with increasing SiCpl and exhibited a sharp decrease when the SiCpl was 15 vol% compared with 5% and 10% SiC addition. At the incipient stage of oxida tion  of  the  SiCpl/ZrB2 between mass change and  composites,  the  correlation  time was  linear  in  that  contact  reaction controlled the oxidation process;  and  as  the  time went on,  the  curve  turned to be approxi mately parabolic because the speed of oxygen diffusion  took control of the oxidation process. The main product  of  the oxidation was m-ZrO2, SiO2 (aluminosilicate). During oxidation of  and borosilicate  the ZrB2-based ceramics, volatilization of B2O3 provided a driving force for the hollow quadrangular prismy ZrO2 crystals the hollow qua growth. The longitudinal direction of  drangular prismy ZrO2 crystals was recognized as [002]. The cross-sectional analysis showed that a larger SiC  content  yielded a more  oxidation-resistant  for  SiCpl/ ZrB2 composites, indicated by the decrease in the thickness of the oxidation layer with increasing SiC content.  References  1.  L. Kaufman, E. Clougher,  and J. B. Berkowit,  ‘‘Oxidation Characteristics  of Hafnium and Zirconium Diboride,’’  J. Trans. Metal. Soc. AIME, 239  458-458 (1967).  2.  F. Monteverde and A. Bellosi,  ‘‘Development and Characterization of Metal Diboride-Based Composites Toughened with Ultra-Fine SiC Particulates,’’  Solid State Sci., 7 622-630 (2005).  3.  W. G.  Fahrenholtz,  ‘‘Thermodynamic Analysis  of ZrB2-SiC Oxidation: Soc., 90 143-148  Formation of  a SiC-Depleted Region,’’  J. Am. Ceram.  (2007).  4.  J. W. Zimmermann, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Thermophysical  Properties of ZrB2 and ZrB2-SiC Ceramics,’’ J. Am. Ceram. Soc., 91 1405- 1411 (2008).  5.  S. M. Zhu, W. G. Fahrenholtz,  and G. E. Hilmas,  ‘‘Influence of Silicon  Carbide Particle Size on the Microstructure  and Mechanical Properties of  Zirconium Diboride-Silicon Carbide Ceramics,’’  J. Eur. Ceram.  Soc., 27  2077-2083 (2007).  6.  L Weng, X. H. Zhang, and J. C. Han,  ‘‘The Effect of Si3N4 on Microstructure, Mechanical Properties and Oxidation Resistance of HfB2-Based Composite,’’ J. Compos. Mater, 43 113-123 (2009).  7.  W. B. Han, P. Hu, and X. H. Zhang, ‘‘High-Temperature Oxidation at 1900  degrees C of ZrB2-xSiC Ultrahigh-Temperature Ceramic Composites,’’ J. Am. Ceram. Soc., 91 3328-3334 (2008).  8.  F. Y. Yang, X. H. Zhang, and J. C. Han,  ‘‘Ablation Mechanism of ZrB2-SiC and C-sf/ZrB2-SiC Ultra-High Temperature Ceramic Composites,’’ J. Inorg. Mater., 23 734-738 (2008).  9.  T.  Zhu,  ‘‘Densification, Microstructure  and Mechanical  Properties  of  ZrB2-SiCw Ceramic Composites,’’ (2009).  J.  Eur. Ceram.  Soc.,  13  2893-2901  10. F. Peng,  ‘‘Oxidation Resistance of Fully Dense ZrB2 with SiC, TaB2, and TaSi2 Additives,’’ J. Am. Ceram. Soc., 91 1489-1494 (2008). S. F. Tang, ‘‘Fabrication and Characterization of Ultra-High-Temperature  11.  Carbon Fiber-Reinforced ZrB2-SiC Matrix Composite,’’ J. Am. Ceram. Soc, 90 3320-3322 (2007).  12. C. Kaya and F. Kaya,  ‘‘Processing, Toughness Improvement and Microstruc tural Analysis of SiC Platelet-Reinforced Al2O3/Y-TZP Nano-Ceramic Matrix Composites,’’ Mater. Sci. Eng., A247 75-80 (1998).  13. C. Kaya and F. Kaya,  ‘‘On the Toughening Mechanisms of SiC Platelet-Re inforced Al2O3/Y-TZP Nano-Ceramic Matrix Composites,’’ Ceram. Int., 25 359-366 (1999).  14. B.-J. Choi and Y.-H. Koh, ‘‘Mechanical Properties of Si3N4-SiC Three-Layer Composites Materials,’’ J. Am. Ceram. Soc, 81 2725-2728 (1998).  15. M. F. Wang, C.-A. Wang, and X. H. Zhang,  ‘‘Preparation and Mechanical  Properties of Zirconium Diboride Matrix Composite Reinforced by Silicon  Carbide Platelets,’’ Rare Met. Mater. Eng., 38 913-915 (2009).  16. F. Monteverde  and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramic in Dry Air,’’ J. Electrochem. Soc., 150 B552-B559 (2003).  17. Z. Xinghong, H. Ping,  and H.  Jiecai,  ‘‘Structure Evolution of ZrB2-SiC During the Oxidation in Air,’’ J. Mater. Res., V23 1961-1972 (2008).  18. H. Jiecai, H. Ping, and Z. Xinghong,  ‘‘Oxidation-Resistant ZrB2-SiC Composites at 2200,’’ Compos. Sci. Technol., v3-4 799-806 (2008).  19. F. Monteverde and A. Bellosi,  ‘‘The Resistance to Oxidation of HfB2-SiC Composite,’’ J. Eur. Ceram. Soc., 25 1025-1031 (2005).  20.  S. N. Karlsdottir,  ‘‘Oxidation of ZrB2-SiC: Solid and Liquid Oxide Phase Formation,’’ J. Am. Ceram. Soc., 92 481-486  Influence of SiC Content on  (2009).  www.ceramics.org/ACT  Oxidation Behavior of ZrB2 Ceramic Matrix Composites  185  \\x0c']"
},{
  "_id": 143,
  "PDF": "Oxidation Behavior of Tantalum Boride Ceramics.pdf",
  "Text": "['Solid State Phenomena Vols 124-126 (2007) pp 819-822 © (2007) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.124-126.819  Online: 2007-06-15  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, www.ttp.net. (ID: 130.216.129.208, University of Auckland, Auckland, New Zealand-01/04/15,04:51:23)  Oxidation Behavior of Tantalum Boride Ceramics Jun-ichi Matsushita1,a, Geum Chan Hwang1 and Kwang Bo Shim2 1Department of Materials Science, Tokai University, 1117 Kitakaname, Hiratsuka 259-1292, Japan 2Department of Ceramic Engineering, Hanyang University, Seoul 133-791, Korea aCorresponding author: jmatsu@keyaki.cc.u-tokai.ac.jp   Keywords: Tantalum Boride, Tantalum oxide, Oxidation, High temperature Abstract. The oxidation behavior of tantalum diboride (TaB2) powder at high temperature was investigated in order to determine the possibility of the use of advanced high temperature structural materials. Unfortunately, monolithic TaB2 were known to be chemical stability up to high temperatures. To date, there have been few reports regarding the properties of TaB2 ceramics. The samples were oxidized at room temperature to 1273 K for 5 minutes to 25 hours in air. The weight changes were measured to estimate the oxidation resistance. The oxidation of samples oxidized for short oxidation time of 5 minutes started at 873 K, and the weight gain increased with increasing oxidation temperature. On the other hand, at the oxidation time of above 1 hour, a maximum weight gain value at 973 to 1073 K was observed. However, even if the oxidation temperature was increased an additional weight change slightly occurred. The weight gain of the sample oxidized at 1273 K for 5 minutes to 25 hours was about 40 to 20 % of the theoretical oxidation mass change. According to the powder X-ray diffraction date, the oxidized TaB2 sample was changed to Ta2O5 at 873 K. Finally, the TaB2 showed a good oxidation resistance at high temperature, because the surface film of tantalum oxide (Ta2O5) formed by oxidation acted as an oxidation resistant layer. Introduction Several tantalum boride phases such as TaB, TaB2, Ta2B, Ta3B2, and Ta3B4 have been registered in the X-ray cards of the International Center for Diffraction Data. Among them, tantalum diboride (TaB2) has been shown to be a potentially useful material because of its excellent chemical stability at high temperature. Then, TaB2 has been attracting a great deal of attention as a material for engineering applications at high temperature, because of its favorable properties such as high melting point (3343 K), chemical stability, and good electrical conductivity [1-3]. Ito et al. reported that a TaB2 had good hardness [1]. To date, there have been few reports regarding the oxidation resistance properties of TaB2 ceramics. In the present study, the isothermal oxidation of a TaB2 powder is investigated from room temperature to high temperatures in order to determine its suitability for advanced high-temperature applications. Experimental Procedure For the TaB2 powder, we used commercial material made by New Japan Metal Co., Ltd., Japan (average particle size of 3 µm and purity of 99 %, one batch for all tests). Isothermal high-temperature oxidation of the TaB2 powder in air was carried out in an electric furnace maintained from room temperature to 1273 K for 25 hours. The TaB2 as-received sample was put on a bed made of a high purity alumina sintered substrate in the hot zone of an electric furnace at operating temperature. After the prescribed period, the sample was removed from the furnace, and allowed to cool. Then, the mass changes were measured to estimate the oxidation resistance using an electromagnetic balance. A few typical oxidized samples were subjected to X-ray diffraction (XRD) study for phase characterization. The diffraction peaks were identified using the computer program of the International Center for \\x0c', '820  Advances in Nanomaterials and Processing  Diffraction Data (ICDD). The oxidized surfaces of the samples were observed using a scanning electron microscope (SEM) to estimate the microstructures. The thermal analysis of TaB2 powder was measured by Thermogravimetry / Differential Thermal Analysis (TG-DTA). Result and Discussion In general, the evaluation of oxidation resistance of B4C or SiC at high temperature is carried out by investigating the weight gain accompanying the following reactions [4, 5]:  B4C (s) + 7/2 O2 (g) \\x07 2 B2O3 (s) + CO (g) \\x01\\x01\\x01 (1) SiC (s) + 3/2 O2 (g) \\x07 SiO2 (s) + CO (g) \\x01\\x01\\x01 (2)  Then, TaB2 is thought to form the tantalum oxide and boron oxide (B2O3) [6]. Even if the weight change is apparently by the vaporization of B2O3 formed as one of the oxidation products at high temperature, the weight gain of the sample should increase with the passage of oxidation time. Thus, the weight gain via oxidation of the sample was measured in order to evaluate the relative oxidation resistance.\\x03The samples were oxidized at room temperature to 1273 K for 25 hours in air. The weight changes were measured to estimate the oxidation resistance. The oxidative weight gain of the TaB2 fine powder versus oxidation time at high temperature are shown in Fig. 1. The oxidation of samples oxidized for short oxidation time of 5 minutes started at 873 K, and the weight gain increased with increasing oxidation temperature. On the other hand, at the oxidation time of above 1 hour, a maximum weight gain value at 973 to 1073 K was observed. However, even if the oxidation temperature was increased an additional weight change slightly occurred. The weight gain of the sample oxidized at 1273 K for 5 minutes to 25 hours was about 40 to 20 % of the theoretical oxidation mass change. Based on the results of the X-ray diffraction analysis, Ta2O5 peaks were present on the oxidized surface of the monolithic TaB2 powder oxidized at 873 K for 1 hour.  \\x0e\\r\\r\\x0e\\r\\x0f\\r\\x10\\r\\x11\\r\\x12\\r\\x13\\r\\x11\\r\\r\\x12\\r\\r\\x13\\r\\r\\x01\\r\\r\\x01\\r\\r\\x01\\r\\r\\x0e\\r\\r\\r \\x05\\x04/,9\\x0443\\x03902507,9:70\\x03(cid:26)\\x0c(cid:28)\\x030\\x04\\x04\\x049\\x03\\x04,\\x043\\x03(cid:10)\\x08(cid:12)\\x122\\x043\\x10\\r2\\x043\\x0e\\x04\\x0e\\r\\x04\\x0f\\x12\\x04  Fig. 1. Weight gain with oxidation as a function of temperature and time.  SEM micrographs of the surfaces of the TaB2 powder oxidized from 673 to 1173 K for 25 hours are shown in Fig. 2. The surface oxidized at 673 K showed a structure almost identical to the powder oxidized for 25 hours. It was found that a coarsely grained oxide layer was formed on the surface of the sample oxidized at 1073 K for 25 hours. \\x0c', 'Solid State Phenomena Vols. 124-126  821   a) As-received sample                   b) Oxidized at 673 K for 25 h       c) Oxidided at 873 K for 25h           d) Oxidided at 1073 K for 25 h     e) Oxidided at 1273 K for 25 h  Fig. 2. SEM images of the oxidized TaB2 powder samples.  Figure 3 shows TG-DTA curves of the TaB2 powder. Oxidation was started with TaB2 powder from about 873 K. Based on the results of the X-ray diffraction analysis, silicon oxide and boron oxide was postulated to be present on the surface of the samples oxidized from about 873 K, and Ta2O5 peaks were postulated to be present on the oxidized surface of the sample. The TaB2 samples showed a good oxidation resistance at high temperature, because the oxidized surface film of Ta2O5 and B2O3 formed by oxidation acted as an oxidation resistant layer. The samples showed a good oxidation resistance at   Fig. 3. TG-DTA curves of the TaB2 powder.  \\x0c', '822  Advances in Nanomaterials and Processing   a) As-received b) Oxidized at 673 K for 25 h c) Oxidized at 773 K for 25 h d) Oxidized at 873 K for 25 h    : TaB2   : Ta2O5                       Fig. 4. X-ray diffraction patterns of oxidized TaB2 powder. high temperature, because the surface film of oxide formed by oxidation acted as an oxidation resistant layer [6]. Then, TaB2 oxidized above 673 K is thought to form the following oxides:  2TaB2 + 11/2 O2 → Ta2O5 + 2B2O3 \\x01\\x01\\x01 (3) Conclusions The oxidation behavior of tantalum diboride (TaB2) powder at high temperature was investigated in order to determine the possibility of the use of advanced high temperature structural materials. Unfortunately, monolithic TaB2 were known to be chemical stability up to high temperatures. To date, there have been few reports regarding the properties of TaB2 ceramics. The samples were oxidized at room temperature to 1273 K for 5 minutes to 25 hours in air. The oxidation of samples oxidized for short oxidation time of 5 minutes started at 873 K, and the weight gain increased with increasing oxidation temperature. On the other hand, at the oxidation time of above 1 hour, a maximum weight gain value at 973 to 1073 K was observed. In conclusion, the TaB2 showed a good oxidation resistance at high temperature, because the surface film of tantalum oxide (Ta2O5) formed by oxidation acted as an oxidation resistant layer. Acknowledgements The author would like to gratefully thank Prof. Koichi Niihara, Extreme Energy-Density Research Institute, Nagaoka University of Technology for helpful advice on boride materials. We would like to thanks Mr. Hiroshi Kiyuna, Mr. Hisayuki Yura, Mr. Yuki Furumura, and Mr. Kengo Ishii of Department of Materials Science, Tokai University for research support. References [1] H. Ito, Y. Satoh, S. Kodama and S. Naka: J. Ceram. Soc. Jpn. Vol. 98, No. 3 (1990), p. 264.  [2] S. Otani, M. M. Korsukava and T. Mitsuhashi: J. Crystal Growth Vol. 194 (1998), p. 430. [3] A. Evstigneeva, R. Singh, M. Trenary and S. Otani: Surface Sci. Vol. 542 (2003), p. 221. [4] J. Matsushita, H. Nagashima and H. Saito: J. Ceram. Soc. Jpn. Vol. 98 (1990), p. 1172. [5] T. Pelekh and J. Matsushita, J. Ceram. Soc. Jpn: Vol. 110 (2002), p. 228. [6] J. Matsushita and S. Komarneni, Materials Research Bull: Vol. 36 (2001), p. 1083.   3040506070d)c)a)b) intensity (a.u.)2θ/o  (Cu Kα)\\x0c', 'Advances in Nanomaterials and Processing   10.4028/www.scientific.net/SSP.124-126   Oxidation Behavior of Tantalum Boride Ceramics   10.4028/www.scientific.net/SSP.124-126.819   DOI References  [6] J. Matsushita and S. Komarneni, Materials Research Bull: Vol. 36 (2001), p. 1083.  doi:10.1016/S0025-5408(01)00560-8         \\x0c']"
},{
  "_id": 144,
  "PDF": "Oxidation behavior of vacuum plasma-sprayed hafnium–tantalum nitrides.pdf",
  "Text": "['REVIEW  This section of Journal of Materials Research is reserved for papers that are reviews of  literature in a given area.  Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  Bradford C. Schulz  Department of Metallurgical & Materials Engineering, The University of Alabama, Tuscaloosa, Alabama 35487-0202, USA  Daniel Butts  Plasma Processes, LLC, Huntsville, Alabama 35811, USA  Gregory B. Thompsona)  Department of Metallurgical & Materials Engineering, The University of Alabama, Tuscaloosa, Alabama 35487-0202, USA  (Received 16 February 2015; accepted 15 June 2015)  A series of (HfN)1\\x00x(TaN)x, ceramics with x representing the starting powder blend compositions of 0.0, 18.8, 28.1, and 46.7 at.%, have been fabricated by vacuum plasma spraying. During the plasma spraying, the mixture lost approximately 25 at.% nitrogen facilitating the precipitation of metallic and metal-rich nitride phases. These specimens underwent static air oxidation exposure up to 1700 °C. In general, it was found that the addition of tantalum nitrides to the hafnium nitrides resulted in poorer oxidation behavior. However, the 18.8 at.% specimen deviated from this trend and had the lowest observed mass change. This specimen formed a dark-colored oxide scale, indexed as Hf6Ta2O17, which acted as a passivation layer. Within the scale, hafnium oxynitride phases were observed. A transformation pathway in forming these rhombohedral oxynitride phases is proposed by the filling in of oxygen in the light element interstitial locations of the rhombohedral e-Hf3N2 and f-Hf4N3 structures.  I.  INTRODUCTION  Ultra high-temperature ceramics (UHTC) comprise a class of materials characterized1-3 by high-melting points , chemical inertness, high hardness,4-7 and moderate oxidation resistance.8-16 These ceramics typically comprise either group IVB or VB transition metals bonded with boron, carbon, or nitrogen. Based on the M:X ratio, where M is the meta l species and X is the interst itial l ight element (boron, carbon, or nitrogen), various metal-rich M6X5, M4X3, M3X2, and/or M2X compounds can be stabilized. The properties of these UHTCs can be further engineered to achieve specific performance metrics by tailoring the phase types and fractions within the microstructure . In this work, the oxidation behavior of hafnium nitrides with various tantalum nitride al loying additions has been investigated. To the authors’ knowledge, there have not been any studies addressing how tantalum nitrides may alter the oxidation behavior of the hafnium nitrides.  Contributing Editor: Tania Paskova a)Address all correspondence to this author. e-mail: gthompson@eng.ua.edu DOI: 10.1557/jmr.2015.191  Oxidation in ceramics commonly occurs by matrix diffusion-controlled9,10,12,13,17 and/or boundary-controlled processes.11,17-19 In diffusion-controlled conditions, the oxidation rate can be regulated by the diffusion rate of oxygen or metal through the oxide product. In a boundarycontrolled process, cracking within the microstructure provides fast-track pathways for the oxygen to penetrate deeper into the material and accelerate oxidation. The hafnium nitrides have been reported to exhibit both matrix diffusionand boundary-controlled oxidation mechanisms.18,20 During this oxidation process, an oxide scale typically forms. The low plasticity of the oxide scale can hinder the relaxation of the scale’s growth stresses. This leads to microcracking and decohesion of the scale-termed spallation.17 Hence tailoring of phases and compositions that can lead to the formation of an adherent, protective oxide scale is advantageous in that it can reduce the matrix and/or boundary oxidation.18 Hafnium oxide has three polymorphs, which are the (a) monoclinic a-HfO2 following: (P21/c), which forms tetragonal b-HfO2 below 1743 °C, (b) that cubic d-HfO2 forms above 1743 °C, and (c) which is stable above 2200 °C up to the melting point.21,22 If the oxide scale forms on the nitride matrix and/or goes through a series of temperatures where one of  (P42/nmc) Fm\\x163m  ð  Þ,  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  Ó Materials Research Society 2015  2949  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  these phase transitions occurs, the volumetric change will generate additional stress states that can promote decohesion. Desmaison-Brut and Montintin18 reported that the oxidation rate for hafnium nitride was lower between 800 and 1000 °C as compared to higher temperature exposures because of a possible intermediate oxynitride phase that had similar kinetic behavior as the monoclinic a-HfO2. Their work was carried out in a carbon dioxide environment. To date, little work has been done exploring the structure or composition of this oxynitride phase or how tantalum nitride additions may influence its formation. In this research, hafnium nitride with various amounts of tantalum nitride, processed by vacuum plasma spray deposition, was studied. The oxidation behavior is addressed as a function of phase content and the formation of complex oxide scale products.  II. EXPERIMENTAL  A. Materials and methods  Four specimens of (HfN)1\\x00x(TaN)x, where x is based on the starting blended powder compositions of 0.0, 18.8, 28.1, and 46.7 at.%, were fabricated by vacuum plasma spraying (VPS). The specimens are referenced based on the ratio of the starting blend powders. Postspraying, the ceramics were chemically analyzed by a light element combustion analysis technique, commonly referred to as LECO testing, to determine any compositional changes that could occur during VPS. The VPS process flows powders into a plasma torch where they are melted and form a rapidly solidified deposit on a substrate surface, which for these materials was a graphite mandrel. Hence, the molten powders and resulting deposit can change composition due to vapor pressure differences between the powders’ constituents and yield a final phase content different than the initial phase content. The starting blend powders, d-HfN and d-TaN, were sieved with a 325 mesh. These powders were synthesized via a self-propagating high temperature synthesis (SHS) reaction of tantalum and hafnium in the presence of nitrogen. The initial tantalum and hafnium had a purity of 99.5 and 99.2%, respectively. After SHS, the d-TaN and d-HfN were determined to contain 46.7 and 49.0% at.% nitrogen, respectively, with less than 800 ppm of oxygen. A custom-constructed VPS stainless steel chamber was evacuated to 10 Pa using mechanical pumps, whereupon it was backfilled with argon to a pressure of 26.6 kPa.23 A volume mixture of 70:1 Ar:H2 served as the plasma gas for the feedstock powder through the VPS tungsten cathode plasma gun operated at a processing power of 45 kW. This gas composition has been shown to increase the processing temperature and lower porosity in deposited material.23,24 The as-sprayed material was deposited to an approximate thickness of 20 mm.  ically sectioned into 250 mm \\x02 200 mm \\x02 20 mm test Post-VPS fabrication, the nitride material was mechanspecimens for the characterization and oxidation studies.  B. Specimen characterization  The characterization of the nitrides consisted of optical and electron beam imaging as well as electron and x-ray diffraction (XRD) analysis for the preoxidation and postoxidation conditions. A FEI F20 Tecnai (scanning) transmission electron microscope (STEM), operated at 200 keV, was used to provide both microstructure imaging and chemical analysis using electron dispersive x-ray spectroscopy (EDS). For the microstructure imaging, high-angle annular dark field (HAADF) was used, as imaging,25 it is based on atomic number or Z-contrast which readily reveals chemical partitioning between the phases, i.e., hafnium and/or tantalum-rich phases will appear bright because of their higher atomic number than a light element—oxygen or nitrogen-phases. The TEM specimens were prepared using a FEI Quanta 3D Dual electron beam-focus ion beam (FIB) microscope that permitted site specific extraction and thinning of the desired TEM foils.26,27 Scanning electron microscopy (SEM) imaging of the initial powders as well as the preoxidation and postoxidation surfaces was performed in a JOEL 7000 SEM operated at 20 keV. This microscope was equipped with the Oxford Instruments electron backscattered diffraction (EBSD) platform, which provided phase identification within the microstructure images. The SEM specimens were mounted and mechanically polished using a 3 lm diamond paste with a further polish for 24 h in aqueous 0.05 lm silica slurry using a Vibromet.26,27 The phase content was identified by XRD using a Bruker D8 Discovery General Area Diffraction Detector System operated with Cu Ka radiation as the source at a setting of 45 keV and 40 mA. All XRD scans were taken from the exposed top surface of the bulk of the specimen. The collected peaks were compared with the phase data found in the International Center for Diffraction Data (ICDD), given in Refs. 28-48. Prior to XRD analysis, the preoxidation specimens were crushed and ground using a mortar and pestle, which ensured a random texture and sufficiently large sampling volume from the powder particulates. Postoxidation XRD analysis was conducted directly on the scale formed on the surface. An XRD line scan was also conducted with a 0.8 mm spot size with 0.01 mm steps for dwell times of 180 s. The reported accuracy of significant figures is based on the 2h step size.  C. Specimen oxidation  In the oxidation experiment, both the preand postoxidation specimen masses were measured. The samples  2950  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  B. Oxidation treatment  In the oxidation experiments, an increase in maximum temperature led to an increase in mass change, Table II. Preand postoxidation optical images of the specimens are shown in Figure 2, with magnified images for clearer viewing provided to the right of the image. Comparing Table II and Fig. 2, as the mass change increased, there was a corresponding presence of an oxide scale, evident by the change in color of the specimens and prevalent powder (which is the delaminated scale) around the bulk of the specimen. The only exception with respect to the spalling was the 18.8 at.% specimen, which had a darkcolored scale as compared to the white or brown color observed in the other specimens. The oxide scale phases for all the postoxidized specimens are shown in the XRD spectra in Figs. 3(a)-3(d). Because of the multiphase nature of the specimens and the various oxide possibilities, there are several overlapping diffracted peaks. This is evident by considering each phase’s line indicators shown at the bottom of each XRD scan in Fig. 3; the reader is referred to Fig. 1’s XRD line indicators for the metallic and metallic nitride phases. Consequently, phase identification was challenging.  were cut to be approximately 12 mm \\x02 12 mm \\x02 12 mm. Since slight sample size variations could exist among them, the mass change for each composition is reported as a percentage change with respect to its initial mass allowing a direct comparison among them. The specimens were placed in silica boats and heated under atmospheric conditions in a 99.5% alumina tube inserted in a Sentro Tech Corp 1700 high-temperature tube furnace with MoSi2 heating elements. The specimens were heated from room temperature to either 700, 1200, or 1700 °C. The 700 and 1200 °C maximum temperature experiments consisted of heating to 200 °C at a 20 °C/min, where it was held for 10 min, then ramped up to the targeted maximum set point at a rate of 10 °C/min, where it was held for an hour and then allowed to cool to room temperature, which took approximately 2 h. The 1700 °C maximum temperature experiments were the same up to 1200 °C but followed with a 10 min hold at 1200 °C and then ramped to 1700 °C at the same rate, where it was held for one hour and allowed to furnace cool afterward, which took approximately 3 h. This heating procedure was the furnaces manufacturer’s recommendations to achieve the higher temperature set points.  III. RESULTS  A. Preoxidation  The average init ial powder sizes were between 6-7 lm for both d-HfN and d-TaN powders. The as-sprayed specimen porosity and nitrogen content is tabulated in Table I. The porosity was measured using the Delesse Principle, which compares the surface area fraction of the total pore count divided by the total surface area of the image.49 The nitrogen measurements were provided by the light element combust ion analysis technique, commonly referred to as LECO testing . The XRD scans for the four preoxidized nitride specimens confirmed the phases of a-Hf, b-Ta, d-HfN, d-TaN, e-Hf3N2, f-Hf4N3, and c-Ta2N, Figure 1. The identification of the metal and metal-rich nitrides confirmed the loss of nitrogen, as noted in the LECO measurements.  TABLE I. Analysis of the four (HfN)1\\x00x(TaN)x, where x is the starting blend composition of 0.0, 18.8, 28.1, and 46.7 at.%. Post-VPS processed data of the porosity and nitrogen content.  Targeted starting chemistry (HfN)1\\x00x(TaN)x  (at.%)  Postprocessed  Porosity %  Nitrogen (at.%)  0.0  18.8  28.1  46.7  2.78  5.58  2.52  3.67  21.16  24.04  17.80  25.70  FIG. 1. Preoxidation normalized XRD scans for the (HfN)1\\x00x(TaN)x specimens, where x is the starting blend 0.0, 18.8, 28.1, and 46.7 at.%. The various phase indicators are presented below for the four specimens with the Greek annotated symbols for the identified phases placed above the experimental XRD peaks.  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  2951  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  In cases of peak overlap, we have complimented our XRD phase identification using the visual cues from the oxide scale color in deducing the dominate phase, i.e., a white scale supported the presence of HfO2 verses a brown scale suggesting a dominate Ta2O5 phase. This does not necessarily discredit that other phases, where peak overlap occurs, are not present, but it does help in determining the most likely or dominate phase fraction. To assist the reader, we have highlighted the most dominate oxide phases to each experimental XRD scan with a faint-colored box—the yellow denotes a-HfO2, light blue Hf6Ta2O17, and light green a-Ta2O5. Those specimens with dominant white-colored scale/ powder had strong a-HfO2 peaks located at 28.36° and 31.68° 2h. The specimens that developed a browncolored scale, notably the 28.1 and 46.7 at.% heated to 1700 °C, have a high-intensity a-Ta2O5 peak at 35.5° 2h. This is a clear case of the need for cross correlation to the visual oxide scale color. One may note that hafnia also has peaks near 35° 2h, but in Fig. 2, a brown oxide scale is observed indicative that Ta2O5 is present even though the hafnia peak at 28.4° 2h (no overlap with tantalum oxide) can be seen. Similar to hafnium oxides, tantalum oxides have different polymorphs at different temperatures. These are the orthorhombic L-Ta2O5 (Pmm2)  TABLE II. Postoxidation mass nitrides, where is the starting and 46.7 at.%.  x  change for blend amount  the of  (HfN)1\\x00x(TaN)x 0.0, 18.8, 28.1,  Specimen chemistry (HfN)1\\x00x(TaN)x  (at.%)  Postoxidation Mass change %  700 °C  1200 °C  1700 °C  0.0  18.8  28.1  46.7  0.37  0.04  1.67  1.35  0.75  0.02  3.34  3.25  1.61  0.10  N/A  3.61  ð  P\\x161  Þ (stable between (below 700 °C),50 triclinic b-Ta2O5 700 °C and 1320 °C), and monoclinic a-Ta2O5 (C2/m) (1320 °C till its melting temperature at 1872 °C)51 phases. The 18.8 at.% specimen, which had a dark-colored scale, Fig. 2, had XRD peaks—33.7° and 34.4° 2h—that could be indexed to the orthorhombic Hf6Ta2O17 phase48 at each temperature condition, highlighted by the faint b lue box . Though the peak near 35° 2h a lso over laps the hafn ia and tan ta lum oxides , the co lor of the sca le sugges ts tha t this peak is no t from those phases . The dark sca le color for the 18 .8 a t .% spec imen , as we l l as the lack of exper imen ta l d iffrac t ion peaks for e i ther tan ta lum ox ide be tween ;25 to ;30° 2h, hafn ia or adds fur ther ev idence tha t the sca le is Hf6Ta2O17 . A t the presence of a-HfO2, a-Ta2O5, became more 1700 °C, readily apparent.  IV. DISCUSSION  The post-VPS process (HfN)1\\x00x(TaN)x specimens’ combustion analysis, Table I, indicated that the nitrogen content was depleted from the initial blend mixture upon spraying. In the case for the Hf:N ratio, as nitrogen is lost, the metal-rich rhombohedral f-Hf4N3 and e-Hf3N2 phases were stabilized; similarly, as nitrogen was lost in the Ta:N ratio, a hexagonal c-Ta2N phase precipitated. In addition, all the specimens showed some fraction of elemental hafnium and/or tantalum. As a consequence of this multiphase microstructure, a complex oxidation behavior occurred. The control 0.0 at.% specimen (no TaN addition) had moderate mass changes as compared to either the 28.1 or 46.7 at.% specimens, Table II. The XRD scale (28.4° 2h) of this specimen was indexed as a-HfO2, confirming that the white powder was hafnia. In contrast, the 28.1 and 46.7 at.% had a strong primary b-Ta peak in the  FIG. 2. Preand postoxidation optical images of the (HfN)1\\x00x(TaN)x, where x is starting blend composition of 0.0, 18.8, 28.1, and 46.7 at.%. Note the dark scale and lack of oxide powder for the 18.8 at.% specimen compared to the other specimens. Enlarged images for each specimen from the 1700 °C/1 h are provided for easier viewing.  2952  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  FIG. 3. Postoxidation normalized XRD scans for the (HfN)1\\x00x(TaN)x, where x is the starting blend composition of (a) 0.0, (b) 18.8, (c) 28.1, and (d) 46.7 at.%. The possible oxide phase identification information is placed below each scan. Specifically, faint color boxes are provided over the experimental and phase indicator data to help and guide the reader to the dominate oxide phases identified. The yellow denotes a-HfO2, light blue Hf6Ta2O17, and light green a-Ta2O5.  postsprayed condition. Since tantalum52 oxidizes faster than hafnium53 and was prevalent in the material, this explains the observed poorer oxidation resistance. The XRD scans of these scales were primarily a-HfO2 and a-Ta2O5. The change in oxide scale resulted in a change in color from white to brown as seen in Fig. 2. The specimen with the lowest mass change, 18.8 at.%, contained a lower b-Ta peak as compared to a-Hf and prevalent close-packed-like structures of e-Hf3N2, f-Hf4N3 and c-Ta2N phases, Fig. 1. The low mass change for this 18.8 at.% specimen was also in agreement with the XRD phase identification that this specimen did not readily oxidize, Table II, or spall, Fig. 2. At 1700 °C, the 28.4° 2h peak for a-HfO2 was evident for this specimen, and the onset of a white scale powder was noted (Fig. 2), which coupled well with a subtle mass change tabulated in Table II. At each oxidation temperature, this specimen had the clearest XRD evidence for the Hf6Ta2O17 phase, Fig. 3. Though the higher tantalum nitride content specimens, 28.1 or 46.7 at.%, also could be indexed with this phase, the scale color change to brown is suggestive that Ta2O5 was dominate. By alloying the hafnium nitride with a smaller amount of tantalum nitride, i.e., 18.8 at.%, the system appears to be able to provide an adequate balance of tantalum to promote the formation of the Hf6Ta2O17 scale without the deleterious effects of Ta2O5 or even HfO2. The complex nature of this orthorhombic oxide scale appears to reduce the rate of oxidation of the nitride. A TEM foil was extracted near the scale-bulk interface, shown in the boxed region labeled (B) for the SEM micrograph in Fig. 4(a), for the 18.8 at.% specimen at 1700 °C. The Hf-rich phase lathes, indicated by the dash arrows, are evident in the STEM-HAADF micrograph in  Fig. 4(b) by their bright (metal-rich) contrast. Figure 4(c) is the EDS line scan (solid arrow) for oxygen content, wh ich inc reases as the probe is scanned over each of the Hf-r ich phases . Th is is sugges t ive tha t these par t icu lar me ta l-r ich n i tr ide phases ac t as oxygen “ge t ters .” As no ted in the SEM m icrograph , F ig . 4 (a) , the ox ide sca le does show numerous microcracks , where oxygen cou ld pene tra te in to the ma ter ia l through the sca le , though cau t ion shou ld be used to deduce if preva len t c rack ing ex is ted a t the e leva ted tempera ture or was a consequence of therma l contrac t ion be tween a l l the phases upon coo l down for pos tmor tem exam ina t ion . These metal-rich rhombohedral f-Hf4N3 and e-Hf3N2based structures54 may provide for the nucleation sites for the prior reported oxynitride Hf7O8N4 phase, which also has the rhombohedral symmetry.42 The f-Hf4N3 structure has one-quarter and the e-Hf3N2 structure has one-third of its light element interstitial sites vacant in the crystal structure, Fig. 4(d). Assuming that the oxygen fills these vacant interstitial sites, the composition of these “transitional” oxynitride phases would be f-Hf4OxN3 and e-Hf3OxN2. Unfortunately, discernable oxynitride phase identification proved difficult in the XRD phase scan due to the similar Hf7O8N4 peak locations near 31° and 35° 2h to the Hf6Ta2O17 phase, Fig. 3. However, by tracking the lattice parameter shift of the f-Hf4N3 and e-Hf3N2 phases, one may deduce if oxygen site filling is occurring by corresponding changes in the lattice parameter. A XRD line scan that transversed from the bulk through the scale was performed with the lattice parameters for the f-Hf4N3 and e-Hf3N2 phases tabulated in Table III. Using these XRD-determined lattice parameters with the proposed “transitional” oxynitride rhombohedral  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  2953  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  FIG. 4. The 1700 °C postoxidation (HfN)1\\x00x(TaN)x x 5 18.8 at.% specimen (a) SEM micrograph revealing the bulk and oxide scale interface with regions of interest for the (S)TEM foil lift out by FIB—denoted as box (B)—and EBSD phase map region box (E). (b) STEM-HAADF micrograph, which dash arrows point to the Hf-enriched nitride lath-like microstructure. The solid arrow indicates the EDS line scan direction. (c) STEM-EDS line profile revealing increases in the oxygen intensity at the Hf-enriched nitride lath positions. (d) Proposed oxygen filling of interstitial vacant sites in the Hf-rich nitride rhombohedral structures. (e) EBSD phase map revealing transitional oxynitride phase formations within the scale.  TABLE III. Lattice parameters (left) based on peak shift from the 18.8 at.% 1700 °C oxidized specimen. The values are compared to prior literature (right most column). The 0.0 mm represents within the bulk of the ceramic moving toward the oxide scale whose outer edge was at 0.50 mm.  Bulk ! Oxide scale  Position  0.00 mm  0.25 mm  0.38 mm  0.50 mm  Literature  Prototype formula  Crystal system  Space group  a (nm)  c (nm)  Prototype formula  Crystal system  Space group  a (nm)  c (nm)  ...  ...  ...  0.3206  2.322  ...  ...  ...  0.3214  3.022  ...  ...  ...  0.3206  2.322  ...  ...  ...  0.3214  3.022  e-Hf3N2 Rhombohedral R\\x163m 166  ð  Þ  0.3206  2.446 f-Hf3N2 Rhombohedral R\\x163m 166 0.3214  ð  Þ  3.044  ...  ...  ...  0.3200  2.467  ...  ...  ...  0.3214  3.075  ...  ...  ...  0.3201 (Ref. 30)  2.326 (Ref. 30)  ...  ...  ...  0.321 (Ref. 31)  3.112 (Ref. 31)  structures—f-Hf4OxN3 and e-Hf3OxN2—along with the previously reported Hf7O8N4 phase,42 an EBSD phase map, Fig. 4(e), was taken from a portion of the scale shown as the box region labeled (E) in the SEM micrograph of Fig. 4(a). The proposed f-Hf4OxN3 and e-Hf3OxN2 phases were successfully identified in the EBSD phase map adding further confidence to their existence. To the authors’ knowledge, these proposed intermediate oxynitride phases of Hf4OxN3 and Hf3OxN2 have not been previously reported. By filling these interstitial vacancies, the oxynitride layer could be providing a further barrier for oxygen diffusion.  If additional oxygen is still required to pass through this phase, it will either require substitutional site exchange and/or a phase transformation to another oxide, such as the Hf6Ta2O17 phase. Further work is still needed to determine the sequence of oxidation; however, the presence of these transitional oxynitrides near the bulk oxide scale interface does suggest that they are likely initial phases that accommodate the onset of oxidation in these hafnium nitrides. How these metallic and metal-rich nitride phases developed in the final-deposited material can be inferred by considering the phase equilibrium within this system  2954  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  with increasing Ta composition. Though VPS is a highly dynamic process and can lead to nonequilibrium phases, phase diagrams can provide very useful insights into how the system may be driven towards equilibrium.23 The approximate location for each post-VPS-sprayed specimen, using the postsprayed LECO nitrogen results, is plotted in the ternary Hf-Ta-N phase diagram55 in Fig. 5. Each specimen’s location is numbered as (1)-(4) on the left side of the phase diagram. Clearly, the experimental presence of a variety of phases, Fig. 1, than those allowed in the phase diagram indicates that the experimental system was not in equilibrium. Regardless, one can still glean how phase fractions may be driven to achieve equilibrium. As the tantalum amount increased, which is moving from (1) to (4) along the Hf-Ta line in Fig. 5, coupled with the nitrogen lost from 50 to 25 at.% (shown by the two dashed red lines), the system is driven further into phase fields where the rhombohedral hafnium-rich nitride e-Hf3N2 and f-Hf4N3 phases are present. This volume fraction increase is confirmed by the XRD intensity increases of these phases relative to a-Hf in Fig. 1. Though these hafnium-rich nitrides existed at higher tantalum nitride additions too, the higher fraction of  metallic b-Ta (e.g., the system is being driven toward the b-Ta phase field in Fig. 5) caused the system to retain a phase that has poorer oxidation properties, which dominated the oxidation at those compositions. The 18.8 at.% specimen appears to be a good balance of stabilizing the rhombohedral hafnium-rich nitride phases for oxidation improvement while maintaining a lower amount of available metallic Hf and Ta to form the bimetallic Hf6Ta2O17 scale. I t is in teres t ing to no te tha t the 18 .8 a t .% spec imen a lso had the h ighes t poros i ty of the spec imens , Tab le I , wh ich one nom ina l ly assumes to y ie ld the wors t ox ida t ion res is tance because such m icros tructure defec ts a l low easy pa thways for ox ida t ion . In con tras t , th is spec imen had super ior ox ida t ion res is tance . Th is ind ica tes tha t the forma t ion of the complex oxide scale of bo th bime ta l l ic and oxyni tride phases is an effec t ive means for ox ida t ion res is tance, even in sys tems wi th less than des irab le m icros tructure conso l ida t ion . Once the 18 .8 a t .% spec imen was exposed to a h igher tempera ture (1700 °C) , this sca le became less effec t ive and a stronger presence of a-HfO2 (ev iden t by the XRD peak iden t ifica t ion in F ig . 3) was observed .  FIG. 5. Ternary Hf-Ta-N phase diagram at 1000 °C based from Rudy.55 Note, the e-Hf2N phase in the original report,54 which was later shown to be incorrect,54 has been replaced by the authors with e-Hf3N2 & f-Hf4N3, which are the correct phases. Numbers (1)-(4) identify the four specimens based on their TaN content of (1) 0.0, (2) 18.8, (3) 28.1, and (4) 46.7 at.%. The dashed lines with arrows point in the direction of nitrogen loss. The red-dashed line indicates the starting and postprocessed nitrogen content. The intersection of these dashed lines is the approximate composition for the four specimens. As the Ta content is increased, and the nitrogen content is reduced (or lost), the system is driven further into the mixed metalrich nitrogen phase fields and toward the metallic Ta phase field.  J. Mater. Res., Vol. 30, No. 19, Oct 14, 2015  2955  \\x0c', 'B.C. Schulz et al.: Oxidation behavior of vacuum plasma-sprayed hafnium-tantalum nitrides  V. CONCLUSION  A series of (HfN)1\\x00x(TaN)x compositionally mixed powders, where x 5 0.0, 18.8, 28.1, and 46.7 at.%, were processed by VPS. The resulting nitrides lost approximately 25 at.% nitrogen during the spraying operation. The loss of nitrogen is associated with differences in vapor pressure between the species under the high temperature and the vacuum environment processing conditions of VPS. This nitrogen loss facilitated the formation of metal-rich nitride f-Hf4N3, e-Hf3N2, c-Ta2N, and elemental metal phases in the final microstructure. The formation of these phases has been explained by the system being driven further into their respective phase fields with an increase in Ta and loss of N using a ternary phase diagram as a guide. Oxidation experiments were conducted between 700 and 1700 °C. The dominate oxide scale composition a-HfO2 a-Ta2O5 when included and the specimen of a-Hf and/or b-Ta contained a significant fraction in the microstructure. The exception was the 18.8 at.% specimen, which formed a dark oxide scale indexed as Hf6Ta2O17 with smaller formations of oxynitride phases within the scale. These collectively reduced the rate of oxidation. Using shifts in the XRD scans, coupled with EBSD phase mapping, transitional f-Hf4OxN3 e-Hf3OxN2 and oxynitride phases were proposed and are believed to be intermediate transitions prior to the stabilization of the reported Hf7O8N4 phase.42 STEM-HAADF and EDS line profiles confirmed that the hafnium-rich nitride phases had increases in oxygen content near the oxide scale with corresponding shifts in the XRD-determined lattice parameters. The presence of f-Hf4N3 and e-Hf3N2 phases within the preoxidized specimen appears to promote the stabilization of these transitional oxynitride phases as well as the ultimate formation of a passivating bimetallic oxide scale that forms from the elemental species.  ACKNOWLEDGMENT  The authors thank Dr. Eric Wuchina for technical discussions. 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},{
  "_id": 145,
  "PDF": "Oxidation Behavior of Zirconium Diboride Nanoparticles.pdf",
  "Text": "['ISSN 0020-1685, Inorganic Materials, 2018, Vol. 54, No. 6, pp. 550-557. © Pleiades Publishing, Ltd., 2018. Original Russian Text © G.V. Kalinnikov, A.A. Vinokurov, S.E. Kravchenko, N.N. Dremova, S.E. Nadkhina, S.P. Shilkin, 2018, published in Neorganicheskie Materialy, 2018, Vol. 54, No. 6, pp. 579-586.  Oxidation Behavior of Zirconium Diboride Nanoparticles  G. V. Kalinnikova, A . A . Vinokurova, S. E. Kravchenkoa, N. N. Dremovaa, S. E. Nadkhinaa, and S. P. Shilkina, *  aInstitute of Problems of Chemical Physics, Russian Academy of Sciences, pr. Akademika Semenova 1, Chernogolovka, Moscow oblast, 142432 Russia *e-mail: ssp@icp.ac.ru  Received September 18, 2017; in f inal form, January 15, 2018  Abstract—The products of oxidation of ZrB2 powders with average particle sizes of ~100 and ~30 nm by atmospheric oxygen under isothermal conditions and during heating have been characterized by thermal analysis, X-ray diffraction, scanning electron microscopy, IR frustrated total internal ref lection spectroscopy, energy dispersive X-ray analysis, and elemental analysis. The oxidation onset has been observed at 594 and 396°C, respectively. Oxidation at temperatures of ≥ 800°C leads to the formation of boron oxide and monoclinic ZrO2, independent of the particle size of ZrB2. The reaction rate constants for the oxidation of ZrB2 nanoparticles ~100 and ~30 nm in size have been determined to be 0.03, 0.15, and 0.31 h-1 at 600, 650, and 700°C and 0.11, 0.35, and 0.81 h-1 at 500, 600, and 700°C, respectively. The apparent activation energies for the oxidation of the ZrB2 nanoparticles ~100 and ~30 nm in size are 161 ± 4 and 62 ± 3 kJ/mol, respectively, as evaluated from the temperature dependence of the rate constants at the above temperatures.  Keywords: zirconium diboride nanoparticles, oxidation, phase, isotherm, rate constant, apparent activation energy for the oxidation process  DOI: 10.1134/S0020168518060067  INTRODUCTION  Group IV metal diborides are used as structural and functional materials  in a variety of  industrial applications and are among potentially attractive components of protective coatings on articles intended for operation at high  temperatures,  in aggressive media, or under accelerated wear conditions [1]. The use of boride materials in a nanostructured state allows one to envisage a considerable extension of their application f ield and stimulates studies of specif ic features of their interaction with corrosive media [2]. Since a nanostructured state of matter has an inherently nonequilibrium nature, nanomaterials are extremely sensitive to high temperatures, aggressive media, deformation loads, etc. Under such conditions,  a nanostructure may undergo  irreversible changes, naturally losing its unique physical, chemical, mechanical, and other properties peculiar  to nanomaterials [3, 4]. In connection with this, understanding general relationships  in  the  behavior  of nanoparticulate borides in aggressive media is a priority issue. Zirconium diboride is a typical Group IV metal diboride. The oxidation of ZrB2 compacts, micron particle size powders, and various composites based on them, has been the subject of extensive studies, which focused on  not only thermodynamic aspects and kinetic correlations of the oxidation process but also on specif ic features of the phases resulting from oxidation (see, for example, Refs. [5-14]). Analysis of available data on the oxidation of ZrB2 demonstrates  the oxidation kinetics of zirconium diboride powders and the phase composition of the oxidation products are determined in many respects by their history, particle morphology and size, the presence of impurities, etc. Results reported by different groups are often inconsistent with each other and characterize the oxidation of only a particular sample. Moreover, there is very limited information about the effect of particle size on the oxidation behavior of ZrB2 and, to the best of our knowledge, there is no such information for nanoparticulate borides [4]. This circumstance led us to study the oxidation of zirconium diboride nanoparticles.  EXPERIMENTAL  Starting reagents. ZrB2 powders with a particle size of ~30 nm were prepared by reacting ZrCl4 and NaBH4. Zirconium diboride powders with a particle size of ~100 nm were synthesized by reacting commercially available boron powder and zirconium powder  550  \\x0c', 'INORGANIC MATERIALS    Vol. 54    No. 6    2018  OXIDATION BEHAVIOR OF ZIRCONIUM DIBORIDE NANOPARTICLES  551  prepared  by decomposing  zirconium hydride  as described by Fokin et al. [15, 16]. The reaction was run in a Na2B4O7 ionic melt as described elsewhere [17-19]. The average particle size dav (Table 1) was evaluated from electron microscopy data and specif ic surface area S measurements under the assumption that the particles were spherical  in shape, using  the wellknown formula d = 6/S γ , where γ  is density. The crystallite size Dhkl was estimated from X-ray powder diffraction data by the Scherrer formula: Dhkl = kλ /(β cos θ )  (along the normal to the hkl planes), where k is the anisotropy coeff icient (taken to be 0.9 for the hexagonal zirconium diboride lattice), λ  is the X-ray wavelength ( λ  = 1.54178 Å for CuK α ), θ  is the diffraction angle, and β  is the full width at half maximum of the diffraction peak.  Electron micrographs of the starting ZrB2 powders are presented in Fig. 1. It follows from these data that the particles are rather uniform in size and nearly spherical in shape.  Table 1. Characteristic temperatures of the ZrB2 particles oxidized during heating  dav is the average particle diameter evaluated from electron microscopy data; d is the particle diameter evaluated from S measurements.  dav, nm  Dhkl, nm   d, nm  Temperature, °С  oxidation onset  exotherm I  exotherm II  ~100  ~85  ~90  594  676  725  ~30  ~28  ~32  396  656  720  Fig. 1. (a, b) Electron micrographs and (c, d) X-ray diffraction patterns of the starting ZrB2 powders (a, c) 100 and (b, d) 30 nm in particle size.  120  100  80  60 2θ, deg  40  20  0  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  500  1000  (b)  100 nm  100 nm  (a)  (c)  60 2θ, deg (d)  001  100  101  102 110 111 002  200  201  112  202 103 211 210  120  100  80  40  20  0  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  500  1000  1500  001  100  101  102 110 111 002  201 112 200  103 202 211 210          \\x0c', '552  KALINNIKOV et al.  Table   2. Temperature-dependent  energy of reaction (1)  enthalpy   and Gibbs  Т, К  Δ Н, kJ/mol  S, J/(mol K)  Δ G, kJ/mol  673  773  873  973  1073  -20 41.4  -2010.8  -2006.1  -2001.2  -1996.5  -437.6  -396.1  -389.8  -384.1  -379.2  -1746.9  -170 4.6  -1665.8  -1627.5  -1589.6  According to elemental analysis and energy-dispersive X-ray spectroscopy data, the composition of the  synthesized  zirconium  diboride  powders  is ZrB2.01-2.05, and they crystallize in hexagonal symmetry (Fig. 1) with lattice parameters a = 0.3168-0.3174 nm and с = 0.3528-0.3535 nm, in reasonable agreement with previously reported data for ZrB2: a = 0.3169 nm and с = 0.3533 nm [20]. Experimental procedure. The oxidation of the ZrB2 nanoparticles in air during heating was studied using a  (a)  DSC, mW/mg  676°C  725°C  Δm = 26.60%  TG, % 145 140 135 130 125 120 115 110 105 100  100  TG, %  130  125  120  115  110  105  100  200  300 400 500 600 700 Temperature, °C (b)  Δm = 18.28%  800  900  DSC, mW/mg  656°C  Δm = 16.06%  720°C  Δm = 17.29%  100  200  300 400 500 600 700 Temperature, °C  800  900  Fig. 2. Thermal analysis results on the oxidation of the ZrB2 particles (a) ~100 and (b) ~30 nm in size.  Netzsch STA 409 PC Luxx thermoanalytical system in combination with a QMS 403 C Aёolos quadrupole mass spectrometer. The process was run at a constant heating rate of 10°C/min in f lowing air with a f low rate of 100 mL/min in the temperature range from 20 to 1000°C. The ZrB2 particles were oxidized with atmospheric air under isothermal conditions at different temperatures in a tubular quartz reactor 20 mm in diameter and 300 mm in length (150-mm-long heating zone). Samples were placed in a platinum foil boat. The temperature in the reactor was maintained with an accuracy of ±2°C by a PT200 temperature controller. The temperature was monitored with an F266 digital device, using a Chromel-Alumel thermocouple as a sensor. The maximum holding time of the samples at preset temperatures was 8 h. The air f low rate in the reactor was 30 mL/min. The degree of conversion, α , was evaluated as the ratio of the measured weight gain of the ZrB2 sample in a given period of time to the maximum possible weight gain calculated for the following ZrB2 oxidation reaction:  ZrB2 + 2.5O2   t   ZrO2 + B2O3.  (1)  Characterization techniques. The phase composi tion of the samples was determined by X-ray diffraction on an ADP-2 diffractometer (monochromatized CuK α  radiation). The  samples were  characterized  by  electron microscopy and energy dispersive X-ray (EDX) spectroscopy on a Zeiss Supra 25 f ield emission scanning electron microscope (SEM) equipped with an INCA x-sight X-ray spectrometer system. Electron microscopic images were obtained at low electron beam accelerating voltages, ~4 kV. EDX analysis data were collected at an accelerating voltage of ~8 kV. The composition of the air used to oxidize ZrB2 was monitored using an MI-1201V mass spectrometer. IR frustrated total internal ref lection (FTIR) spectra were measured in the range from 500 to 4000 cm-1 using a PerkinElmer Spectrum 100 Fourier transform spectrometer and a Vertex 70V  spectrometer, both equipped with accessories for taking ref lection spectra. The specif ic surface area (S) of the samples was determined using a Quadrasorb SI analyzer. Nitrogen and oxygen in solid phase were determined by EDX analysis, as well as on an Elementar Vario Micro cube CHNS/O element analyzer. Zirconium was determined by chelatometric titration in the presence of Xylenol Orange. Boron was determined by potentiometric titration of the mannitol-boric acid complex with an alkali after zirconium precipitation from the analyte solution and by energy dispersive X-ray analysis using standard procedures.  INORGANIC MATERIALS    Vol. 54    No. 6    2018  18 16 14 12 10 8 6 4 2 0  18 16 14 12 10 8 6 4 2 0  ⎯ ⎯ → Δ \\x0c', 'INORGANIC MATERIALS    Vol. 54    No. 6    2018  OXIDATION BEHAVIOR OF ZIRCONIUM DIBORIDE NANOPARTICLES  553  RESULTS AND DISCUSSION  In the temperature range 400-800°C, reaction (1) has a high thermodynamic feasibility and yields zirconium dioxide and boron oxide in a condensed state (Table 2). The formation of zirconium borate at these temperatures was left out of consideration [21]. The Gibbs energy values obtained in this study agree well with those reported by Poilov and Pryamilova [14]. In our calculations, we used reference data from Chase [22].  Figure 2 presents thermal analysis results on the oxidation of the ZrB2 particles ~100 and ~30 nm in size during heating from 20 to 1000°C. It is seen from these data that, independent of the particle size, the DSC curves each have two exothermic peaks, whose temperatures are indicated in Table 1 together with the oxidation onset temperatures. It follows from Table 1 that the particle size has a particularly strong effect on the oxidation onset temperature and the position of exotherm I. By analogy with the oxidation of TiB2  powders of various particle sizes, the considerable difference in oxidation onset temperature can be tentatively attributed to the role of the deformation induced in the particles of various sizes by the Laplace pressure (P = 2 σ /r, where σ  is surface tension and r is the particle radius). The deformation may act to accelerate the oxidation of smaller ZrB2 nanoparticles at lower temperatures [13].  Figures 3a-3e show X-ray diffraction patterns of the products of isothermal oxidation of the ZrB2 powder with a particle size of ~100 nm at temperatures of 500, 600, 650, 700, 800, and 1000°C (heating to 1000°C). At a temperature of 500°C, the phase composition of the sample remained unchanged and there was no weight gain. At 600°C, ZrO2(T) and ZrO2(M) begin to form (Figs. 2a, 3b). At 650°C (Fig. 3b), two zirconium dioxide polymorphs were also observed to form: tetragonal ZrO2 (T) and monoclinic ZrO2 (M). At 700°C, larger amounts of these phases were formed  Fig. 3. X-ray diffraction patterns of the products of isothermal oxidation of the ZrB2 powder with a particle size of ~100 nm at temperatures of (a) 600, (b) 650, (c) 700, (d) 800, and (e) 1000°C (heating to 1000°C) and the ZrB2 powder with a particle size of ~30 nm at temperatures of (f) 400, (g) 500, (h) 600, and (i) 700°C.  100  120  60 80 2θ, deg  40  20  400  200  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  100  120  60 80 2θ, deg  40  20  0  600  200  400  800  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  100  120  80  40 60 2θ, deg  20  0  1000  500  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  0  (a)  1500  500  1000  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  0  2000  1500  1000  500  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  600  400  200  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  0  ZrB2 ZrO2(M) ZrO2(T)  1000  1500  2000  500  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  0  ZrB2 ZrO2(M) ZrO2(T)  500  250  750  1000  1250  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  0  ZrB2  600  400  800  200  I  n  t  e  n  s  i  t  y  ,  a  r  b  .  u  n  i  t  ZrB2 ZrO2(T)  ZrO2(T)  ZrB2 ZrO2(T)  ZrB2 ZrO2(M)  ZrO2(M)  ZrB2 ZrO2(M)  (b)  (c)  (d)  (e)  (f)  (g)  (h)  (i)                                      \\x0c', '554  KALINNIKOV et al.  tetragonal phase ZrO2(T) transforms into the monoclinic phase ZrO2(M):  ZrO2(T)   t   ZrO2(M).  (2)  Under these conditions, the starting crystalline phase ZrB2 completely oxidizes according to scheme (1).  Figure 4 shows SEM images of the products of oxidation of the ZrB2 nanoparticles with a particle size of ~100 nm at 700 and 800°C and with a particle size of ~30 nm at 700°C. There are well-def ined changes in nanoparticle morphology and size relative to the original state (Fig. 1), due to reactions (1) and (2).  The f irst exothermic peaks in the DSC curves for the oxidation of ZrB2 with particle sizes of ~100 and ~30 nm at temperatures of 676 and 656°C (Fig. 2) are attributable to the onset of ZrO2 and B2O3 formation according to scheme (1), even though no B2O3 phase was detected in the X-ray diffraction patterns of the samples oxidized at different  temperatures (Fig. 3). According to the EDX analysis data in Fig. 5, there is an excess oxygen concentration (relative to the 26 wt % O in the stoichiometric composition of ZrO2), which may be due to the formation of boron oxide or boric acid. The IR spectroscopy results in Fig. 6 support the former assumption. The IR FTIR spectrum of the ZrB2 oxidation products is essentially identical to that of individual boric anhydride [23], without vibrations characteristic of H3BO3 (at 3200, 1450, or 1196 cm-1 [24]).  The second exothermic peaks in Fig. 2, due to the oxidation of the ZrB2 particles ~100 and ~30 nm in size in the temperature range 720-725°C, are presumably attributable to further formation of B2O3 and the phase transition (2), as evidenced by the X-ray diffraction results in Figs. 3c, 3d, 3h, 3i.  According to X-ray diffraction, EDX analysis, and chemical analysis data, the oxidation products contained no nitrogen-containing derivatives of zirconium or boron.  Complete oxidation of ZrB2 powder by reaction (1) should be accompanied by a weight gain of 71%. Under the conditions of this study, the weight change was considerably smaller (Figs. 2, 7) and, according to the X-ray diffraction data for the ~100-nm particles, the oxidation products contained the starting ZrB2 phase (Figs. 3b-3e). The reduced weight gain might be due, f irst, to the volatility of boron oxide (its melting and boiling points are ~450 and ~2250°C, respectively). According  to mass spectrometry data,  the released gas contained a noticeable amount of B2O3 800°C. Another possible even at temperatures of  cause was the formation of a B2O3 glass f ilm on the surface of the ZrB2 particles, which prevented further  INORGANIC MATERIALS    Vol. 54    No. 6    2018  (a)  (b)  (c)  100 nm  100 nm  100 nm  Fig. 4. Electron micrographs of the products of oxidation of the ZrB2 powders with a particle size of ~100 nm at 700 (a) and 800°C (b) and with a particle size of ~30 nm at 700°C (c).  (Fig. 3c). Starting at 800°C and at higher temperatures, the tetragonal phase disappeared and there was only the monoclinic phase of ZrO2 (Figs. 3d, 3e).  The ZrB2 powder with a particle size of ~30 nm showed a different type of oxidation behavior. X-ray diffraction patterns of the products of isothermal oxidation of the ZrB2 powder with a particle size of ~30 nm at temperatures of 400, 500, 600, and 700°C are presented in Figs. 3f-3i. It follows from these data that, in the temperature range 400-500°C, the oxidation process is accompanied by the formation of only tetragonal zirconium oxide, ZrO2(T). The process reaches completion at 600°C. At 700°C,  the metastable,  ≥ ⎯ ⎯ → \\x0c', 'OXIDATION BEHAVIOR OF ZIRCONIUM DIBORIDE NANOPARTICLES  555  Spectrum 1  B  O  0  Element  B K O K Zr L  Total  1 mm  Spectrum 1  Zr  Zr  2  Arb.  conc. 6.73 32.99 33.85  Intensity  corr. 0.7528 1.5857 0.8337  B  4  Weight %  12.71 29.58 57.72  100.00  6 8 Energy, keV  Weight %  sigma 2.89 1.90 2.68  Atomic %  32.14 50.56 17.30  Spectrum 1  (a)  (b)  Spectrum 1  60 μm  O  0  Arb.  conc. 3.25 4.48 20.43  Intensity  corr. 0.6427 0.7385 0.9057  Element  B K O K Zr L  Total  Zr  2  Zr  Weight %  15.02 18.02 66.96  100.00  4 6 Energy, keV  Weight %  sigma 6.14 3.15 5.54  Atomic %  42.75 34.66 22.59  Fig. 5. EDX analysis data for the products of isothermal oxidation of the ZrB2 particles (a) ~100 nm in size at a temperature of 800°C and (b) ~30 nm in size at a temperature of 700°C.  2 1 3 6 3 2 1 1  9 8 0 1 8 7  e c  n  a  b  r  o  s  b  A  4000  3500 3000 2500 2000 1500 Wavenumber, cm-1  1000  500  Fig. 6. IR FTIR spectrum of the products of isothermal oxidation of the ZrB2 powder with a particle size of ~100 nm at 700°C.  diffusion of oxygen atoms and zirconium diboride oxidation reaction.  Figure 7 shows kinetic curves for the oxidation of ZrB2 powder particles ~100 and ~30 nm in size at  various temperatures. The kinetic curves of the ZrB2 samples with a particle size of ~100 nm are well represented by an Avrami-Erofeev equation of the form [-ln(1 - α )]1/n = kτ  (where α  is the degree of conver INORGANIC MATERIALS    Vol. 54    No. 6    2018  \\x0c', '556  KALINNIKOV et al.  α  0.7  0.6  0.5  0.4  0.3  0.2  0.1  0  (a)  800°C  700°C  650°C  α  600°C  2  4 Time, h  6  8  0.6  0.5  0.4  0.3  0.2  0.1  0  (b)  700°C  600°C  500°C  400°C  2  4 Time, h  6  8  Fig. 7. Degree of conversion α temperatures.   as a function of oxidation time for the ZrB2 particles (a) ~100 and (b) ~30 nm in size at different  sion and τ  is time), with n = 1/2, a value characteristic of vapor-solid heterogeneous processes. The reaction rate constants were determined to be 0.03, 0.15, and 0.31 h-1 at 600, 650, and 700°C, respectively. The apparent activation energy for the oxidation of the ZrB2 nanoparticles is 161 ± 4 kJ/mol as evaluated from the temperature dependences of the rate constants at 600, 650, and 700°C.  The kinetic curves of the ZrB2 sample with a particle size of ~30 nm can also be described by an Avrami-Erofeev equation, but, unlike in the case of the oxidation of the ZrB2 powder with a particle size of ~100 nm, the parameter n is 1/3 rather than 1/2, suggesting an increase in the diffusion contribution to the kinetics of the oxidation process. The reaction rate constants were determined to be 0.11, 0.35, and 0.81 h-1 at 500, 600, and 700°C, respectively. The apparent activation energy  for  the oxidation of  the ZrB2 nanoparticles is 62 ± 3 kJ/mol as evaluated from the temperature dependence of the rate constants in the range 500-700°C, which is considerably lower than the 161 ± 4 kJ/mol obtained for the ZrB2 samples with a particle size of ~100 nm.  Thus, in the range of oxidation temperatures studied here, the particle size of ZrB2 has no effect on the phase composition of the f inal oxidation products, but inf luences predominantly the oxidation rate and the temperature at which the oxidation products appear. Complete oxidation of ZrB2 according to scheme (1) was only observed at a particle size of ~30 nm.  ACKNOWLEDGMENTS  This work was supported by the Russian Federation Ministry of Education and Science, state research tar get, theme no. 0089-2014-0028, state registration no. 01 201 361 876.  REFERENCES  1.  boridy   Vysokotemperaturnye   Serebryakova, T.I., Neronov, V.A., and Peshev, P.D., (High-Temperature Borides), Chelyabinsk: Metallurgiya, 1991. 2. Andrievskii, R.A., The role of nanoscale effects in the interaction between nanostructured materials and envi ronments, Prot. Met. Phys. Chem. Surf., 2013, vol. 49,  no. 5, pp. 528-540. 3. Andrievski, R.A. and Khatchoyan, A.V., Nanomaterials  5.  in Extreme Environments, Fundamentals and Applica tions, Cham: Springer International, 2016. 4. Andrievskii, R.A., Nanostructured  titanium, zirconium, and hafnium diborides: synthesis, properties, size effects, and stability, Usp. Khim., 2015, vol. 84, no. 5, pp. 540-554. Irwing, R.J. and Worsley, J.G., Oxidation of titanium diboride and zirconium diboride at high temperatures, J. Less.-Common. Met., 1968, vol. 16, no. 2, pp. 103- 112. 6. Lavrenko, V.A., Glebov, L.A., and Lugovskaya, E.S., High-temperature oxidation of zirconium diboride in air, Zh. Fiz. Khim., 1973, vol. 47, no. 4, pp. 887-891. 7. Voitovich, R.F. and Pugach, E.A., High-temperature oxidation of Group IV metal borides, Poroshk. Metall. (Kiev), 1974, no. 2, pp. 57-62. 8. Lebugle A. and Montel G., Étude comparée de l’oxidation des diborures de Zr, Hf et Ti, Rev. Int. High Temp. Refract., 1974, vol. 11, pp. 231-244. 9. Ban’kovskaya, I.B., Pevzner, B.Z., and Gorbatova, G.N., Specif ic features of oxidation of powder boron-containing  vol. 72, no. 6, pp. 937-941. 10. Ming Guo, Guo-Jun Zhang, Yan-Mei Kan, and PeiLing Wang, Oxidation of ZrB2 powder in the temperature range of 650-800°C, J. Alloys Compd., 2009, vol. 471, pp. 502-506.  formulations, Russ. J. Appl. Chem., 1999,  INORGANIC MATERIALS    Vol. 54    No. 6    2018  \\x0c', 'OXIDATION BEHAVIOR OF ZIRCONIUM DIBORIDE NANOPARTICLES  557  nanoparticles as a result of reaction between zirconium tetrachloride and sodium borohydride, Inorg. Mater., 2017, vol. 53, no. 8, pp. 80 4-808.  19. Kravchenko S.E., Burlakova A.G., Korobov  I.I., Kalinnikov G.V., Domashnev, I.A., Shilkin, S.P., and Andrievskii, R.A., Synthesis of nanosized group IV borides in ionic melts of anhydrous sodium tetraborate, Russ. J. Inorg. Chem., 2016, vol. 61, no. 4. pp. 429-433.  20. Diagrammy sostoyaniya dvoinykh metallicheskikh sistem:  Spravochnik (Phase Diagrams of Binary Metallic Systems: A Handbook), Lyakishev, N.P., Ed., Moscow: Mashinostroenie, 1996.  21. Kolovertnov, D.V., Oxidation of glass-ceramic composites based on zirconium boride and silicon-containing  (Chem.) Dissertation, St. Petersburg: Grebenshchikov Inst. of Silicate Chemistry, Russ. Acad. Sci., 2012.  compounds, Extended Abstract   of Cand. Sci.  22. Chase, M.W.,  Jr., NIST-JANAF  thermochemical tables, fourth edition, J. Phys. Chem. Ref. Data, 1998, monograph 9.  23. Sidorov, T.A. and Sobolev, N.N., Infrared and Raman spectra of boric anhydride: III. Interpretation of the vibrational spectrum of boric anhydride and calculation of the isotopic effect, Opt. Spektrosk., 1958, vol. 4, no. 1, pp. 9-16.  24. Bethell, D.E. and Sheppard, N., The infrared spectrum and structure of boric acid, Trans. Faraday Soc., 1955, vol. 51, pp. 9-15.  Translated by O. Tsarev  11. Ortiz, A.L., Zamora, V., Rodríguez-Rojas, F., A study of the oxidation of ZrB2 powders during high-energy ball-milling in air, Ceram. Int., 2012, vol. 38, pp. 2857- 2863. 12. Kolovertnov D.V., Ban’kovskaya I.B., Yuritsyn N.S. Thermogravimetric investigation of the oxidation of the ZrB2-SiO2 composite in the temperature range 800- 1300°C, Glass Phys. Chem., 2008, vol. 34, no. 4, pp. 461-469. 13. Andrievskii, R.A., Shul’ga, Yu.M., Volkova, L.S., Korobov,  I.I., Dremova, N.N., Kabachkov, E.N., Kalinnikov, G.V., and Shilkin, S.P., Oxidation behavior of TiB2 microand nanoparticles, Inorg. Mater., 2016, vol. 52, no. 7, pp. 686-693. 14. Poilov, V.Z. and Pryamilova, E.N., Thermodynamics of oxidation of zirconium and hafnium borides, Russ. J. Inorg. Chem., 2016, vol. 61, no. 1, pp. 55-58. 15. Fokin, V.N., Fokina, E.E., and Shilkin, S.P., Synthesis of coarsely crystalline metal hydrides, Russ. J. Gen. Chem., 1996, vol. 66, no. 8, pp. 1210-1212. 16. Fokin, V.N., Fokina, E.E., Tarasov, B.P., and Shilkin, S.P., Synthesis of  the  tetragonal  titanium dihydride  in ultradispersed state, Int. J. Hydrogen Energy, 1999, vol. 24, nos. 2-3, pp. 111-114. 17. Burlakova A.G., Kravchenko S.E., Domashnev I.A., Vinokurov A.A., Nadkhina, S.E., Volkova, L.S., and Shilkin, S.P., Special features of preparation of nanosized zirconium diboride powders of various dispersity, Russ. J. Gen. Chem., 2017, vol. 87, no. 5, pp. 906-911. 18. Kravchenko, S.E., Burlakova, A.G., Domashnev, I.A., Nadkhina, S.E., Dremova, N.N., Vinokurov, A.A., and Shilkin, S.P., Formation of zirconium diboride  INORGANIC MATERIALS    Vol. 54    No. 6    2018  \\x0c']"
},{
  "_id": 146,
  "PDF": "Oxidation Behavior of Zirconium Diboride Silicon Carbide Produced by the Spark Plasma Sintering Method.pdf",
  "Text": "['Oxidation Behavior of Zirconium Diboride Silicon Carbide Produced  by the Spark Plasma Sintering Method  Carmen M. Carney,*,w,z,y  Pavel Mogilvesky,*,z,y  and Triplicane A. Parthasarathy*,z,y  zAir Force Research Laboratory, Materials and Manufacturing Directorate, AFRL/RX  y  UES Inc., Dayton, Ohio 45432  Dense samples of ZrB2-20 vol% SiC were successfully fabricated by spark plasma sintering without the use of sintering aids.  Oxidation behavior of these samples was characterized by exposing them to 14001, 15001, and 16001C in an ambient atmosphere for 150 min, and by measuring the weight gains of the  sample and crucible, as well as the thickness of the oxide scale  and the glassy outer layer. The effects of gravity on the viscous  outer  layer  are  shown  to  result  in  significant  heterogeneity  within a sample. The oxidation scales were  characterized by  scanning electron microscopy and transmission electron micros copy with energy dispersive spectroscopy analysis. The oxide scale was found to be composed of three layers: (1) a SiO2-rich glassy outer layer, (2) an intermediate layer of a ZrO2 matrix with interpenetrating SiO2, and (3) a layer containing a ZrO2 matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions.  I.  Introduction  BECAUSE of their high melting temperatures and excellent resistance to oxidation and evaporative erosion, transition metal diborides (MB2) such as ZrB2 and HfB2, commonly referred to as ultra-high-temperature ceramics (UHTCs), are be ing considered as principal candidates for the leading edges of sharp-bodied reentry vehicles.1-6 ZrB2-based UHTCs have received a majority of the attention due to their lower density (6.09 vs 10.5 g/cm3  for HfB2) and lower cost. Several shown that the addition of SiC improves the oxidation resistance of the diborides at temperatures 412001C.7-12 In terms of  studies have  mechanisms,  it has been shown that ZrB2 oxidizes to ZrO2 and liquid B2O3, which evaporates at higher temperatures (412001C) as B2O3 (g).13-15 The addition of SiC allows the formation of SiO2 (melting temperature 17101C), which is more resistant to evaporation and has a higher viscosity at elevated  temperatures than B2O3. However, B2O3 continues to flux the silica scale, lowering its viscosity.  Processing of diboride UHTCs typically involves the use of (above 19001C)  elevated temperatures  for extended periods of  time (30 min or longer) accompanied by applied pressures. Some  success  has  been met with  pressureless  sintering  of  these  UHTCs, but this process normally requires a sintering aid such as MoSi2 or B4C.16,17 Recently, the spark plasma sintering (SPS) method has provided a processing technique to rapidly (o 5 min at the sintering temperature) densify ZrB2-based UHTCs. However, many of these studies also use a sintering aid in addition to SiC to ensure full density,18-24 unless reactive synthesis  is used.19 These sintering aids may affect the oxidation proper ties of these materials. In this paper we seek to use SPS to pre pare dense, ZrB2-based UHTCs with SiC and without sintering aids.  These samples will be tested for  their oxidation resistance,  with particular attention paid to unanswered questions that re main in the literature, regarding the oxide scale morphology and  chemistry. These questions include the possible effect of exper imental conditions (geometry of the sample and crucible) on the  fluid flow of the viscous glassy outer layer. Although glass flow  has been suggested to be important to boria flow in monolithic diborides,15 it is not fully explained in the two phase materials  that contain silica forming compounds. Additionally, there are  inconsistencies in the reported work on the phase distribution of  the oxidized scale. Some reports suggest a region of  the oxide  scale that is depleted of SiC due to active oxidation, while others  show no such region among samples tested at similar temperatures.7-13  The objectives of this study are to (1) fabricate dense ZrB2- SiC without sintering aids using SPS, (2) study the effect of ex ternal scale viscosity on oxidation mechanisms, and (3) deter mine the change in oxide scale morphology and composition as  a function of  temperature. We show that SPS can be used to  produce dense ZrB2-SiC without Significant sources of error in measuring weight  the need for  sintering aids.  change  are  shown to arise from the viscosity of the scale at temperatures below 16001C. An analysis of the oxide scale morphology with  emphasis on the glassy scale will highlight the importance of the  viscosity of  this layer. Finally, a compositional analysis of  the  oxide scale is presented.  II.  Experimental Procedure  ZrB2 Advanced Materials, East Providence, RI) were used to mix 80 vol% ZrB2 and 20 vol% SiC (ZrB2-SiC). The b-SiC was a 45- 55 nm (reported) powder with 97.5% purity (o0.15% free Si, o0.15% Cl, o0.75% free C, and o1.25% O). All impurities (C, in the ZrB2 powder were o0.02 wt% Fe, Hf, Ti, Al, and Be) except for hafnium, which was 1.3 wt%. The ZrB2 powder had a measured mean starting size of 9.4 mm. Powders were ball-milled  (Millmaster Chemical, New York, NY) and SiC (Reade  using Si3N4 grinding media in isopropanol for 18 h. The powders were then dried while stirring followed by 18 h of dry milling  with the same Si3N4. The dried powders were sieved using an 80-mesh screen. The weight loss of the Si3N4 grinding media (0.03% lost weight) was 0.14 wt%, based on the total powder  weight.  Nine grams of  the dried powder was loaded into a 20-mm graphite die coated with BN and lined with graphite foil. The  sample was sintered using SPS (FCT Systeme GmbH Model HP  D 25-1, Rauenstein, Germany) with a heating and cooling rate of 901C/min and a maximum temperature of 20001C (achieved  using 5.5 V and 1 kA). The hold time was 5 min. The temper ature was measured by an optical pyrometer focused on the bottom of a borehole in the punch B5 mm from the powder. A  D. Butt—contributing editor  This work was supported in part by USAF Contract # FA8650-04-D-5233.  *Member, The American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: ccarney@ues.com  Manuscript No. 25384. Received October 22, 2008; approved April 3, 2009.  Journal  J. Am. Ceram. Soc., 92 [9] 2046 - 2052 (2009)  DOI: 10.1111/j.1551-2916.2009.03134.x  r 2009 The American Ceramic Society  2046  \\x0c', 'September 2009  Oxidation Behavior of Zirconium Diboride Silicon Carbide  2047  vacuum of at least 1.5 Pa was maintained for the entire heating  cycle. The pulse sequence used for heating was 10 ms on, 5 ms  off, with a single pulse. A uniaxial load of 32 MPa was applied at 12001C and held during the heating cycle. Samples with dimensions of 4 mm \\x02 2.5 mm \\x02 2 mm were cut from the puck with a diamond saw and polished on all sides to a 1 mm finish for  oxidation tests.  ZrB2-SiC specimens were exposed to air at temperatures of 14001, 15001, and 16001C for 150 min. The samples were heated and cooled at 201C/min. A horizontal MoSi2 resistance-heated tube furnace was used to heat the samples. The samples were  placed on a zirconia crucible to avoid interaction with the al umina D-tube used to support  the sample in the furnace. All  samples were placed so that the two sides with the largest surface  area were parallel to the zirconia crucible (Fig. 1). The bottom  side is defined as the side facing the zirconia crucible, while the  top faces the centerline of the furnace. Weights before and after  oxidation were measured using a balance with 0.01 mg preci sion. Both the sample and the zirconia crucible weights were  measured before and after oxidation.  The microstructures of the as-sintered and oxidized samples were studied by polishing the samples to a 1 mm finish. All pol ishing was done using diamond slurries. The oxidized samples  were polished perpendicular to the top and bottom faces. Sam ples were also prepared for microscopy using a focused ion  beam microscope  (FIB: FEI Dual Beam 235, FEI, Hillsbor ough, OR). The microstructure was characterized using scan ning electron microscopy (SEM: Quanta, FEI, Hillsborough,  OR)  and  transmission  electron microscopy  (TEM:  Phillips  CM200, FEI, Hillsborough, OR) along with energy dispersive  spectroscopy (EDS) for chemical analysis.  III.  Results and Analysis  (1)  SPS Processing  The bulk density of the sintered pucks was 5.47 g/cm3, as mea sured by the Archimedes’ test method. Using a rule of mixtures  and assuming the true densities of ZrB2 and SiC to be 6.09 and 3.21 g/cm3, the theoretical density would be 5.51 g/cm3, giving a theoretical density of 499% for the sample. Figure 2 is a mi crograph of the ZrB2-SiC sample preceding oxidation. The microstructure is typical of those presented in previous literature.  The microstructure is fairly regular and no open porosity is ob served, which is in agreement with the density calculations. The  darker phase is SiC and appears  to be uniformly distributed  within the ZrB2 matrix.  (2) Weight Gains from Oxidation  Figure 3 is a plot of a parabolic rate constant (Kw) vs 1/T for the samples held for 150 min between 14001 and 16001C. Kw is the square of the weight gain per surface area divided by time, where  the weight gain was the total  increase in weight of  the sample  and crucible. The plot  shows an increase in weight gain with  temperature as expected. At these temperatures, the weight gain  results from the production of B2O3, SiO2, and ZrO2  through  Fig. 2. Scanning electron micrograph of a polished ZrB2-SiC specimen prepared by spark plasma sintering at 20001C for 5 min. The measured density was 5.47 g/cm3.  the following reactions:  ZrB2 þ 5 2  O2 ðgÞ ! ZrO2 þ B2O3 ðl Þ  SiC þ 3 2  O2 ! SiO2 ðsÞ þ COðgÞ  (1)  (2)  It is expected that the vapor pressure of B2O3 increases considerably as the temperature approaches 14001C.25-27 The thick ness of the oxide scales was measured on the sides and bottom of  each sample. At least seven micrographs were taken for both the  bottom and sides of each sample, with magnifications allowing  the oxide scale to fill the image. Three measurements were taken  from each micrograph. The data presented in Table I are the  average thicknesses measured for the glass layer and the total  oxide scale thickness minus the glass thickness for each sample.  As expected, the glass scale increases with temperature on the  bottom side of the sample. Additionally, the total scale thickness  increases with temperature, but  the layer underneath the scale  only increases slightly. This is because of the protective nature of  the glass scale and the potentially lower oxygen partial pressure,  as discussed in the next section. The glass scale thickness on the  side of  the sample is in contradiction and decreases with tem perature. The reason for the decrease in thickness is the increase  in viscosity of the SiO2-rich glass with temperature. As the temperature increases, gravity works to remove the less viscous glass  from the sample sides. The overall effect of the glass removal is a  less protective glass scale that allows more oxygen penetration  and thus a longer total scale thickness that consists primarily of a ZrO2 matrix. The glass scale thins considerably at 16001C as SiO2 approaches its melting point and its viscosity decreases.27  Fig. 1.  Schematic showing the placement of the ZrB2-SiC samples relative to the zirconia crucible. Samples were placed so that the top of the  sample is defined as the side facing the centerline of the furnace.  Fig. 3.  Change in weight with temperature. Kw 5 (w/A)2/t, where (w/A) is the weight gain normalized to the sample surface area and t is the hold time (150 min). Heating and cooling rates were 201C/min.  \\x0c', '2048  Journal of the American Ceramic Society—Carney et al.  Vol. 92, No. 9  Table I.  Oxide Scale Thickness Measured at Each Test  Temperature  Side Oxide Scale  Bottom Oxide Scale  T (1C)  1600  1500  1400  Glass  16.4 mm 35.8 mm 19.9 mm  ZrO2 layer  70.8 mm 50.1 mm 46.7 mm  Glass  20.6 mm 5.9 mm 4.7 mm  ZrO2 layer  15.4 mm 9.4 mm 10.5 mm  The average of at least 20 measurements was calculated separately for the side  and bottom (defined in Fig. 1) of each sample. The glass scale thickness was mea sured separately from the reminder of the oxide scale.  Additionally, the ZrO2 scale becomes thicker due to a less protective outer layer. The glass layer is less protective due to two  factors:  (1)  increased oxygen diffusivity with temperature, and  (2) its decrease in thickness.  Additional proof of glass flow is realized by measuring the  weight change of the sample and zirconia crucible separately. At 14001 and 15001C, the percent of weight gain attributable to the  sample is 92% and 90%, while only 50% of the weight gain is measured from the sample at 16001C. Experiments conducted in  which the  zirconia crucible alone was heated using the  same  heating profile as the samples showed no weight increase by heating in air up to 16001C. Thus the weight change measured  for the zirconia crucibles used underneath the ZrB2-SiC samples must be a result of glass flow from the sample to the crucible.  Visual  inspection of the crucibles confirms the presence of glass  on the crucible.  (3)  Oxide Scale Morphology  A representative image of an oxidized sample is shown in Fig. 4,  while the inset shows the typical morphology of the top corners.  The top corners have less protective SiO2 due to glass flow from the sharp edges, while the bottom corners probably lose SiO2 due to contact with the zirconia crucible. The oxide scale on the  top (side facing the furnace centerline) and sides is not uniform,  but the scale on the bottom (side facing the zirconia crucible) is  fairly even. Other literature has described an undulating oxide scale.10,28,29 Sample nonhomogeneity is often discussed as a rea son for the observed differences, but there is no indication that  there would be a difference between the surfaces of the sample.  The samples were cut  from the bulk, and after polishing they  were placed in the furnace with no regard to original orientation.  Fig. 4. Scanning electron microscopy (SEM) image of the ZrB2-SiC sample oxidized at 16001C for 150 min. The light gray areas are the oxide  scale. The undulating oxide scale is clearly seen when the sample is tilted 351. Ag paste is used to attach the sample to the SEM stub. The inset  micrograph shows a top corner of the sample. There is extensive oxida tion at all  the SiO2-rich phase at allows oxidation to penetrate further into the bulk.  four corners. The lack of  the corner  Fig. 5. Top view of the oxidized ZrB2-SiC sample surface. The sample was oxidized for 3 h at 16001C. Energy Dispersive Spectroscopy (EDS)  showed that the light gray areas contain Zr and O, while the dark gray  areas contain Si and O.  Fig. 6. (a) Scanning electron microscopy (SEM) micrograph of the oxide scales formed after holding at 14001C for 150 min. Region 1 is the  SiO2-rich glassy outer layer, region 2 is the layer of SiO2 infiltration into the porous oxide, region 3 is a layer containing a ZrO2 matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions, and region 4 is the bulk ZrB2-SiC. (b) Plot of the ratio of the Si energy dispersive spectroscopy (EDS) peak to the Si and O EDS peaks, showing the transition  of the dark phases from SiO2 to Si-O-C inclusions to SiC. (c) Plot of the ratio of the Zr EDS peak to the B and O EDS peaks, showing the tran sition of the light phases from ZrO2 to ZrB2.  \\x0c', 'September 2009  Oxidation Behavior of Zirconium Diboride Silicon Carbide  2049  The continuous glass is much more protective and allows less  oxidation, as was shown in Table I.  It is well accepted that the formation of a borosilicate glass on  top of the ZrB2-SiC limits oxygen diffusion to the bulk material and suppresses further oxidation.30 In addition, it may also limit  the diffusion of gases that form within the compact (B2O3, CO, and CO2). If the pressure of the gaseous products exceeds the ambient pressure and the glass is viscous enough to flow, these gases can create bubbles in the glass scale.31 Observations in our  lab of  the samples during the heating cycle suggest  that  these  bubbles  form and burst during the initial hold time, and the  quantity formed decreases with holding time. Bubbles are not  observed to form during cooling.  The viscosity of a borosilicate glass is expected to decrease  with temperature and with increasing B2O3 content. B2O3 exists as a liquid above approximately 4501C and begins to evaporate above 11001C, but  replenished by diffusion from the  can be  be  to  be  would  bulk. At the temperatures under consideration, the glassy scale SiO2-rich.25,27 The through evaporation of B2O3 in the SiO2-rich scale leads to increased viscosity. As the glass becomes more viscous, bubble  continual  expected  loss  formation  is  inhibited  and  the  product  gases must  diffuse  through the layer. Simultaneously, oxidation is  suppressed by  the protective glassy scale, so fewer gases are being formed in side the scale.  The rupture of these bubbles leaves behind a void where ox ygen can rapidly progress through the porous ZrO2, until surface is covered with SiO2. This process was captured in a SEM image (Fig. 5), showing a bubble that had burst and two  the  distinct regions underneath: exposed ZrO2 and the region that had begun to re-seal with SiO2. No bubbles are observed to form on the bottom of the sample. The absence of bubbles on the  bottom helps to contribute to the uniformity of  the glass and  thus the decreased oxide scale thickness. One possible explana tion of the absence of bubbles on the bottom surface may result  from a decreased oxygen partial pressure in the gap between the  sample and the crucible, which would lead to decreased oxida tion. The decreased oxygen partial pressure can be caused by  oxygen consumption and displacement by the outgassing of CO  and B2O3, produced in reactions (1) and (2), cally trapped.  that could be lo (4)  Oxide Scale Composition  Figure 6(a) is a micrograph showing the oxide scale produced when the ZrB2-SiC sample is oxidized at 14001C for 150 min. The samples oxidized at 15001 and 16001C for 150 min had the same morphology as the sample oxidized at 14001C with longer  oxide scales. EDS was used to measure the chemistry of  indi vidual regions of the sample. The ratios of the maximum peak  heights for the oxygen signal and carbon signal relative to the  maximum peak height of silicon are plotted for the Si-containing  regions (darker features), and the ratios of  the maximum peak  heights  for  the oxygen signal and boron signal relative to the  maximum peak height of zirconium are plotted for the Zr-con taining regions (lighter features), as shown in Figs. 6(b) and (c),  Fig. 7. (a) Scanning electron microscopy (SEM) micrograph of a sample tilted to 521. From this image it  is obvious that  the portion of  the  oxide scale that is made up of a ZrO2 matrix polishes more rapidly than the bulk or the regions containing SiO2. The C-cap is applied before the sample is bombarded by Ga ions to protect the surface. (b) Schematic  showing the focused ion beam (FIB) cutting process. The sample was cut  perpendicular  to a polished face. The Ga ions are used to mill away  material so that the inside of the oxide scale can be observed. (c) SEM  micrograph of the FIB cut oxide scales. The image shows the polished  face (left), FIB cut face (middle), and area of the sample that was milled away (right). The sample was oxidized at 16001C for 150 min. Region 1  is the SiO2-rich glassy outer layer, region 2 is the layer of SiO2 infiltration into the porous oxide, region 3 is the layer containing a ZrO2 matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions, and region 4 is the bulk ZrB2-SiC.  Fig. 8.  Compositional analysis of the reaction zone (region 3) and the bulk ZrB2-SiC (region 4). The secondary electron image (far left) corresponds with the X-ray maps of C, O, Si, and B. The focused ion beam (FIB) cut face is enclosed by two solid parallel lines. The double-headed arrow in the  secondary electron image and the Si and C maps indicates the transition from SiC to C-rich inclusions. The single-headed arrows designate areas with a  high B concentration and low O concentration,  indicating partially oxidized ZrB2.  \\x0c', '2050  Journal of the American Ceramic Society—Carney et al.  Vol. 92, No. 9  respectively. These ratios are used to distinguish SiO2 from SiC and ZrO2 from ZrB2. This analysis suggested four regions that are indicated in the micrograph for the sample oxidized at 14001C:  (1) a SiO2-rich glassy outer (2) a ZrO2 matrix with penetrating SiO2, (3) a reaction zone containing ZrB2/ ZrO2, with inclusions that contained varying amounts of Si, C, B, and O, and (4) bulk ZrB2-SiC. The regions are approximately separated by the solid lines in Fig. 6. When the sample was tilted to 521 for imaging in the SEM  layer,  (Fig. 7(a)), it can be observed that the area made up of the ZrO2 matrix was concave, likely from the more rapid polishing that  occurred in the less dense regions. Because it was possible that  grain pullout was occurring in this region, a cut was made in the  sample using a focused ion beam (FIB) microscope in order to  more precisely determine the morphology and composition of  the oxide scale. Figure 7(b) shows the FIB process. First, a car bon cap (Fig. 7(a)) is applied to the sample to protect the surface  and assure a straight parallel cut without rounding the edge, and  then a focused beam of gallium ions is used to remove material  by sputtering perpendicular to a polished face of the sample. The  resulting sample shows a polished face, a FIB-cut  face, and a  milled face (Fig. 7(c)). Because the FIB cutting did not  induce  the extent of damage (grain pullout) seen from polishing, it was  possible to see that the voids of the ZrO2 were actually filled with material, as opposed to just the rough inclusions seen in the  polished cross section (Fig. 6(a)). Spot EDS analysis was used to  determine the composition of all the dark phases in the image.  The dark features in the bulk region correlate with SiC, while the  dark features within the reaction zone have the same size and  shape as the original SiC but are shown to contain higher con centrations of C and O. Near the outer SiO2-rich layer, the dark features are shown to be mostly Si and O. These four regions  agree with the same regions observed in the polished cross sec tion of  the oxidized sample  (Fig. 6). There  is a morphology  difference between the SiO2, the inclusions in the reaction zone, and the unoxidized SiC that can be seen in the micrographs. FIB cuts were made in samples oxidized at 14001 and 15001C and  were shown to have the same morphology.  Figure 8 shows  the  compositional analysis of  the  reaction  zone (region 3) and bulk ZrB2-SiC (region 4) of the sample. The sample was tilted so that the analysis was done on the FIB cut  face (enclosed by the two parallel  lines). From the analysis,  it is  clear that the bulk SiC grains have a different chemistry than the  inclusions found in the reaction zone. A double-headed arrow  marks the approximate transition from SiC in region 4 and in clusions rich in C in region 3. Interestingly, there is also evidence  of ZrB2 grains that are only partially oxidized, as determined by the regions (indicated with arrows in Fig. 8) in the O and B maps  that have a high concentration of B and no O signal.  TEM analysis of a foil thinned by FIB (Fig. 9(a)) shows the  unoxidized crystalline SiC in the bulk and the amorphous phase  containing Si, C, B, and O in the reaction zone. A small Zr peak  in the qualitative EDS analysis (Fig. 9(b)) of the amorphous Si-,  C-, B-, and O-containing region is attributable to either background signal or a small amount of dissolved ZrO2 in the glass.31 The presence of partially oxidized ZrB2 is confirmed by selected area electron diffraction in TEM. Partially oxidized hexagonal  ZrB2 grains surrounded by monoclinic ZrO2 are found throughout the reaction zone (Fig. 9(c)).  Fig. 9.  (a) Transmission electron microscopy (TEM) micrograph show ing the bulk region containing SiC and the reaction zone containing an  amorphous Si-C-B-O. (b) Energy dispersive spectroscopy (EDS) anal ysis of the amorphous Si-C-B-O. (c) TEM micrograph of partially ox idized ZrB2 that is found throughout the reaction zone. The hexagonal ZrB2 and monoclinic ZrO2 were distinguished with selected area electron diffraction patterns.  IV.  Discussion  Second,  the present work shows that  the flow of  the glassy  outer layer has a significant influence on oxidation kinetics mea First,  the  results  from this work, presented in Fig. 2,  clearly  surements. Furnace oxidation is a common tool  for UHTC  show that a dense microstructure can be obtained by sintering  analysis. Although some authors report undulating scales and  ZrB2-SiC using SPS without any sintering additives. Unlike previous studies,18-24 there were no intentional sintering addi tives; however, it is possible that erosion of the Si3N4 grinding media may contribute as a sintering aid.32 Further experimen tation is needed with different grinding media or mixing meth ods to isolate this effect.  samples  that were disregarded due to adhesion to a crucible,  there is little discussion on how the glass viscosity—and thus its  flow—affects the oxidation data. Because few authors report the  type of crucible used, it is unknown if the weight change of the crucible is being taken into consideration. At 14001 and 15001C,  these differences are expected to be only a few percent, but as  \\x0c', 'evidenced in Fig. 3, the weight change of the zirconia crucible at 16001C due to the flow of the glassy layer is significant. We find  that, due to glass flow, the ZrO2 scale and the SiO2 scale thicknesses (and thus the weight gain) both depend on the orientation  of  the  sample. Because  the  scale  thickness varies around the  sample, as a result of glassy layer variations, the weight change is  representative of a set of samples with different random orien tations to gravity. Additionally, the flow from the corners con tributes to excessive oxidation, skewing the results. The bubbles  that are formed in the glassy scale seem to contribute to the  undulating scale and rapid oxidation underneath burst bubbles.  Controlling the viscosity of the SiO2-rich scale with additives may lead to improved oxidation resistance.33  Third, the present work reveals more detailed aspects of the  oxidation product than those reported in prior work. Inclusions  that contained C, Si, B, and O were observed in the reaction  zone. The FIB cut made perpendicular to the oxidation front  showed that these inclusions are the same size and shape as the  original SiC features.  The observation of Si-C-B-O-containing glass inclusions in  the reaction zone is somewhat unique. Others have reported a  SiC-depleted region in a ZrB2 layer just above the unoxidized bulk in samples heated between 15001 while only Monteverde11,29 and Nickel38 show C-rich inclusions in the same layer at 14501 and 15001C. We have observed amor and 22001C,10,25,31-37  phous Si-C-B-O inclusions throughout the reaction zone at temperatures between 14001 and 16001C. These inclusions have  also been observed in our lab using combinations of two differ ent ZrB2 powders, two different SiC powders, different milling media, and HfB2-based UHTCs. The observation of partially oxidized ZrB2 suggests that not all the ZrB2 grains were able to fully oxidize before the protective SiO2 layer formed. Holds at lower temperatures (before SiO2 formation) different morphologies and lifetimes of the samples. Studies be and subsequent heating cycles may  lead to  ing conducted at different  time intervals will help further elicit  the reaction mechanisms during oxide scale formation.  V.  Summary  We have successfully produced dense ZrB2-20 vol% SiC using the SPS processing technique without the use of any sintering  aids. Oxidation studies conducted at temperatures between 14001 and 16001C showed that the flow of external glassy phase  influenced the results significantly. Using independent measure ments of sample and crucible, and analyzing these along with  scale thicknesses, showed that the effect of viscous flow was most significant at 16001C. In addition, observations of flow and  bubble formation in the SiO2-rich outer glassy layer pointed to the importance of the properties of this layer in oxidation pro tection. Finally,  the oxide scale was shown to consist of  three  layers:  (1) a SiO2-rich glassy outer layer of a ZrO2 matrix with interpenetrating SiO2, and (3) a layer containing a ZrO2 matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions.  layer,  (2) an intermediate  Acknowledgments  We would like to thank Kathleen Sevener and John Halloran for their insight ful discussions related to oxide formation and glass flow and Dr. Michael Cinibulk  for his input and guidance throughout the project.  References  1R. Savino, M. D. S. Fumo, D. Paterna, and M. Serpico, ‘‘Aerothermodynamic  Study of UHTC-Based Thermal Protection Systems,’’ Aerospace Sci. Tech., 9,  151-60 (2005). 2R. Monti, M. D. S. Fumo, and R. Savino, ‘‘Thermal Shielding of Reentry Ve hicle by Ultra-High-Temperature Ceramic Materials,’’ J. Thermophys. Heat Trans fer, 20 [3] 500-6 (2006). 3P. Kolodziej, J. V. Bowles, and C. 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V. der Biest,  ‘‘Synthesis and Microstructural Features of ZrB2-SiC-Based Composites by Reactive Spark Plasma Sintering and Reactive Hot Pressing,’’  Scritpa Mater., 57, 317-20 (2007). 24A. Bellosi, F. Monteverde, and D. Sciti,  ‘‘Fast Densification of Ultra-High Temperature Ceramics by Spark Plasma Sintering,’’ J. Am. Ceram. Soc., 3, 32-40  (2006). 25A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, ‘‘Oxidation of Zirconium Diboride-Silicon Carbide at 15001C at a Low Partial Pressure of Oxygen,’’ J. Am.  Ceram. Soc., 89 [10] 3240-5 (2006). 26W. C. Tripp and H. C. Graham, ‘‘Thermogravimetric Study of the Oxidation of ZrB2 in the Temperature Range of 800-15001C,’’ J. Electrochem. Soc., 118 [7] 1195-9 (1968). 27S. N. Karlsdottir, J. W. Halloran, and C. E. Henderson, ‘‘Convection Patterns  in Liquid Oxide Films on ZrB2-SiC Composites Oxidized at High Temperature,’’ J. Am. Ceram. Soc., 90 [9] 2863-7 (2007). 28F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 29F. Monteverde and L. Scatteia,  in Dry  ‘‘Resistance to Thermal Shock and to Oxida tion of Metal Diborides-SiC Ceramics for Aerospace Application,’’ J. Am. Ceram.  Soc., 90 [4] 1130-8 (2007). 30C. W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of SiC Addition on  the Oxidation of ZrB2,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 31S. N. Karlsdottir and J. W. Halloran, ‘‘Zirconia Transport  by Liquid  Convection During Oxidation of Zirconium Diboride-Silicon Carbide,’’ J. Am.  Ceram. Soc., 91 [1] 272-7 (2008). 32F. Monteverde and A. Bellosi,  ‘‘Effect of the Addition of Silicon Nitride on  Sintering Behaviour and Microstructure of Zirconium Diboride,’’ Scripta Mater.,  46 [3] 223-8 (2002). 33Y. Ogura and T. Morimoto,  ‘‘Mass Spectrometric Study of Oxidation of SiC  in Low Pressure Oxygen,’’ J. Electrochem. Soc., 149 [4] J47-52 (2002). 34I. G. Talmy, J. A. Zaykoski, M. M. Opeka, and S. 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Grigoriev ‘‘The  Kinetics of Oxidation of ZrB2-based UHTC Composites,’’ ECI Conference, Ultra High Temperature Ceramics: Materials for Extreme Environment Applications,  August 3-8, 2008, Lake Tahoe, CA.  &  \\x0c']"
},{
  "_id": 147,
  "PDF": "Oxidation behavior of zirconium diboride–silicon carbide at 1800°C.pdf",
  "Text": "['Scripta Materialia 57 (2007) 825-828  www.elsevier.com/locate/scriptamat  Oxidation behavior of zirconium diboride-silicon carbide at 1800 °C  Jiecai Han, Ping Hu,* Xinghong Zhang and Songhe Meng  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, China  Received 1 June 2007; revised 1 July 2007; accepted 2 July 2007  Available online 6 August 2007  ZrB2-based ultrahigh-temperature ceramics, including 20 and 30 vol.% SiC, have been investigated at 1800 °C under different oxygen partial pressures. ZrB2-30 vol.% SiC exhibited the highest oxidation resistance under high oxygen partial pressure, whereas it displayed the lowest oxidation resistance under low oxygen partial pressure. The thickness of the oxide scale of ZrB2 containing either 20 or 30 vol.% SiC increased with decreasing oxygen partial pressure. Moreover, this trend was intensified by increasing the SiC content.  Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: ZrB2; SiC; Oxidation behavior  Ceramic compounds of refractory borides, carbides and nitrides, such as TaC, ZrB2, ZrC, HfB2, HfC and HfN, belong to a group of materials known as ultrahigh-temperature ceramics (UHTC) [1-4]. Interest in UHTC has increased significantly in recent years because of the drive to produce a thermal protection system and other components for hypersonic aerospace vehicles [5-7]. In addition to high melting temperatures, ZrB2 and HfB2 have a unique combination of chemical stability, high electrical and thermal conductivities, and resistance to erosion/corrosion that makes them suitable for the extreme chemical and thermal environments associated with hypersonic flight, atmospheric re-entry and rocket propulsion [8-11]. Oxidation resistance is a major issue for the use of UHTC. Great efforts have been devoted to this aspect and much progress has been gained recently. The mechanical and physical properties of monolithic zirconium (hafnium) diboride are inappropriate for this specific application due to their low oxidation resistance and strength/fracture toughness. The introduction of second phases (i.e. SiC, MoSi2) has succeeded in improving both the oxidation/ablation resistance and mechanical properties of UHTC. In particular, the addition of Ta compounds has been shown to improve the oxidation resistance of ZrB2-SiC below 1800 °C, whereas it is detrimental to the performance of ZrB2-SiC above  * Corresponding author. Tel./fax: +86 45186402382; e-mail: huping@  hit.edu.cn  1800 °C [12,13]. In accordance to the previously reported results, a diboride matrix composite including only SiC as a second phase behaves as one of the most promising compositions. However, the optimization of the composition and microstructure has not been profoundly studied and the oxidation behavior is not clear, especially at temperatures above 1800 °C. The purpose of this paper is to investigate the oxidation behavior of ZrB2 containing 20 and 30 vol.% SiC at 1800 °C under different oxygen partial pressures. The effects of SiC content and oxygen partial pressure on the oxidation resistance of the material are also discussed. The samples used here for oxidation testing were fabricated from commercial ZrB2 (Northwest Institute for Non-ferrous Metal Research, China) and SiC (Weifang Kaihua Micro-powder Co., Ltd., China) powders. The powder mixtures of ZrB2 + 20 vol.% SiC (ZS1) and ZrB2 + 30 vol.% SiC (ZS2) were ball-milled in ethanol for 8 h using WC media and dried in a rotating evaporator. Milled powder was then uniaxially hot-pressed in a boron nitride coated graphite die at 2000 °C for 60 min under vacuum and 30 MPa of applied pressure. Bulk density and theoretical density were evaluated using the Archimedes method and the rule-of-mixture, respectively. Sample coupons 2.0 · 1.0 · 0.35 cm were cut from the hot-pressed materials, and all surfaces were diamond-polished to a 1 lm finish. Coupons were ultrasonically cleaned successively in detergent, deionized water, acetone and alcohol prior to exposure. Samples were loaded into a slotted ZrO2 refractory brick and then  1359-6462/$ see front matter Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2007.07.009  \\x0c', '826  J. Han et al.  / Scripta Materialia 57 (2007) 825-828  exposed to 60 min oxidation in stagnant air at 1800 °C with oxygen partial pressures of 0.2 and 2 · 10\\x004 atm, using a bottom-loading furnace with zirconia heating elements. The total pressure in the furnace was controlled to 1 · 10\\x003 atm for obtaining an oxygen partial pressure of 2 · 10\\x004 atm. Three samples were used for oxidation tests at each condition and the repeat experiments were also performed under the same conditions. X-Ray diffraction (Rigaku, Dmax-rb) was used to identify oxide phases present after exposure. Scanning electron microscopy and energy-dispersive spectroscopy (FEI Sirion, Holland) were used to characterize the composition and microstructure of the surface and cross-section of the samples after testing. The different oxide layers were also investigated after removing the surface layers by polishing parallel to the original surface. The material removal was monitored using optical microscopy so that the desired region was reached. The bulk densities of the sintered ZS1 and ZS2 billets were 5.41 and 5.22 g cm\\x003, which correspond to relative densities of 98.1 and 99.8%, respectively. The relative density increased as the SiC content increased from 20 to 30% by volume. Figure 1 shows a scanning electron micrograph of the polished surfaces of the ZrB2-SiC UHTCs. The microstructure of the composites is regular, with a mean grain size of about 6 lm, and the residual porosity is very scarce. The SiC particulate, which represents the majority of the dark features in Figure 1, is homogeneously dispersed intergranularly within the ZrB2 matrix and no agglomeration was detected. Exaggerated grain growth of ZrB2 was restricted due to the existence of SiC particles. In addition, the introduction of SiC substantially enhanced densification of ZrB2 during hot-pressing. The explanation for the beneficial effect of SiC on densification lies in the fact that most SiC powder particles are known to have an oxidized surface layer, which promotes formation of liquid phases during hot-pressing, leading to increased densification [14]. Figure 2 shows the macrographs of the oxidized ZrB2based UHTCs containing 20 and 30 vol.% SiC at 1800 °C with oxygen partial pressures of 0.2 and 2 · 10\\x004 atm. The macrographs of the oxidized materials appear similar and the configurational stabilities remained after oxidation. Extensive bubble formation was observed on the all samples after oxidation. The size of the bubble for ZS1 oxidized at the oxygen partial pressure of 0.2 atm was greater than for the others. A large amount of glass was formed on the surface of ZS2 after oxidation at 1800 °C with an oxygen partial pressure of 0.2 atm. The thicknesses of the oxide scale are shown in Table 1. ZS2 exhibited the highest oxidation resistance at 1800 °C under the oxygen partial pressure of 0.2 atm,  Figure 1. SEM of  the polished surfaces of  the ZrB2-SiC UHTCs:  (a)  ZS1 and (b) ZS2.  Figure  2. The macrographs of the oxidized ZrB2-SiC UHTCs 1800 °C for 1 h under different oxygen partial pressures: (a) ZS1 (pO2 = 2 · 10\\x004 atm); 0.2 atm); (b) ZS1 (pO2 (c) ZS2 (pO2 = 0.2 atm); and = 2 · 10\\x004 atm). (d) ZS2 (pO2  at  =  Table 1. Thicknesses of  the  scale of ZrB2-SiC oxidized at 1800 °C  under different oxygen partial pressures  Materials  Oxygen partial  Thickness  Thickness of the  pressure (atm)  of the oxide  SiC-depleted  scale (lm)  scale (lm)  ZrB2-20SiC  ZrB2-20SiC  ZrB2-30SiC  ZrB2-30SiC  0.2 2 · 10\\x004  0.2 2 · 10\\x004  300  310  192  380  40  85  54  220  whereas it displayed the lowest oxidation resistance under the oxygen partial pressure of 2 · 10\\x004 atm. The thickness of the oxide scale formed under the oxygen 2 · 10\\x004 atm (380 lm) was partial pressure of almost twice that formed under the oxygen partial pressure of 0.2 atm (192 lm). However, the oxidation behavior of ZS1 did not vary significantly with oxygen partial pressure, and the thicknesses of the scale for ZS1 were 300 and 310 lm at oxygen partial pressures of 0.2 and 2 · 10\\x004 atm, respectively. It appears that the thickness of the oxide scale does not change greatly with oxygen partial pressure for ZS1. This suggests that the ZrB2 containing 20 vol.% SiC has a similar oxidation rate at different oxygen partial pressures. The micrographs in Figure 3 show some representative microstructural modifications of the ZS2 surface after oxidation at 1800 °C under an oxygen partial pressure of 0.2 atm. The surface was covered by a thick silica glassy layer (Fig. 3(a and b)). Bubbles of different sizes were detected in the outermost layer, as shown in Figure 3(a-c). Aggregated zirconia crystals of various sizes and shapes formed underneath the large bursting bubbles (Fig. 3(d)). The formation of the large bubbles was most likely due to the coalescence of the gaseous products (i.e., CO, SiO) inside the external forming glass. The bubbles tend to burst when the vapor pressure exceeds the ambient pressure. The growth of large bubbles beneath the external scale often causes local spalling which, in turn, noticeably reduces the mechanical and oxidation properties. It should be noted that a large number of small white spherical particles were also detected on the surface which are not reported in previous studies. Analysis by EDS showed that they consist of Si, O and Zr. The EDS results may not be from the sampling of SiO2 under or beside the ZrO2 particles since the size of the spherical particles is far larger than the accuracy of the equipment. The formation of the spherical particles is most likely due to the liquid transport of the SiO2-ZrO2 solution to the surface and then coagula \\x0c', 'J. Han et al.  / Scripta Materialia 57 (2007) 825-828  827  Figure 4. Cross-sectional micrographs of oxidized materials at 1800 °C for 1 h: (a) ZS1 (pO2 = 0.2 atm); (b) ZS2 (pO2 = 0.2 atm); (c) ZS1 (pO2 2 · 10\\x004 atm) and (d) ZS2 (pO2 = 2 · 10\\x004 atm).  =  surface was parallel to the surface of the oxidized specimens. The columnar zirconia showed an equiaxed structure in this direction. The columnar zirconia matrix was surrounded by silica glass, which efficaciously seals grain boundaries and plugs pores, leading to restricted permeability to oxygen. Limited B existed in the columnar zirconia matrix, while no B was detected in the silicate, which is in contrast with previously reports in which almost as much B was present as borosilicate [9,16]. A large number of white spherical particles were also observed around the matrix, which was confirmed as Zr and W oxides. W element was introduced into the sample through the WC milling media. WC was oxidized and then transported, along with ZrO2, by gas and liquid convection and deposited in the oxide scale. The thickness of the oxide scale of both ZS1 and ZS2 increased with decreasing oxygen partial pressure, as shown in Table 1. Moreover, this trend was intensified by an increase in SiC content. The cross-sectional micrographs of ZS1 and ZS2 oxidized at 1800 °C with an oxygen partial pressure of 2 · 10\\x004 atm are shown in Figure 4(c) and (d). The oxide scale of ZS1 was similar to that oxidized at an oxygen partial pressure of 0.2 atm, as can be seen in Figure 4(a and c). However, the oxide layer markedly changed for ZS2 when the oxygen partial pressure shifted from 0.2 to 2 · 10\\x004 atm, and the thickness  Figure  5. SEM and EDS  results  for  the  oxidized ZS1  (1800 °C,  pO2  = 0.2 atm, 1 h).  Figure  3. Surface morphologies and EDS of the scale of the ZS2 specimen after oxidation at 1800 °C for 1 h under an oxygen partial  pressure of 0.2 atm.  tion at the silica rich surface layer. The liquid solution of is \\x189% at 1800 °C, ZrO2 in SiO2 estimated from the phase diagram of the binary ZrO2-SiO2 system [15]. The experimental results as shown in Figure 3(e) are in agreement with the ZrO2-SiO2 diagram. The cross-sectioning and polishing of the samples after furnace testing provided insight into the microstructural details of greatest interest. Figure 4 shows the crosssectional micrographs of oxidized materials at 1800 °C under different oxygen partial pressures. Both ZS1 and ZS2 have four distinct layers: (i) a silica rich outer layer, which is believed to act as a protective scale; (ii) a subscale of crystalline zirconia, containing little silicate; (iii) a zirconium diboride region depleted in SiC (the ‘‘SiC-depleted region’’); and (iv) unaltered material. However, the thickness of the scale for ZS2 was far less than that for ZS1. The surface of the oxidized ZS2 was covered by a thick silica glass layer. The formation of an external silica-based glass layer is very effective in limiting the inward diffusion of oxygen into the inner bulk, thus enhancing the resistance to oxidation. Furthermore, the microstructure of the subscale for ZS1 was different from that for ZS2. The subscale of the ZS1 oxidized at 1800 °C under the oxygen partial pressure of 0.2 atm shows an oriented growth, which was observed in this study. The crystalline zirconia exhibited a columnar structure in the oxidized ZS1 sample, whereas this phenomenon was not observed in ZS2 after oxidation in the same conditions. The oriented growth of the scale was most likely due to the evolution of the gaseous by-products, which promoted the growth of the zirconia parallel to the direction of the discharge of the gas products. Figure 5 shows SEM and EDS results oxidized ZS1 (1800 °C, = 0.2 atm). The  for the polished  pO2  \\x0c', '828  J. Han et al.  / Scripta Materialia 57 (2007) 825-828  of the outmost silica glass was remarkably lower than that oxidized at an oxygen partial pressure of 0.2 atm. The reduction of the formed silica glass would be responsible for the increasing oxidation rate. The thicknesses of the SiC-depleted layer increased with decreasing oxygen partial pressure for both ZS1 and ZS2, as shown in Table 1. Apparently, the low oxygen partial pressure is favorable for the active oxidation of SiC. Thus, the thickness of the SiC-depletion layer at low oxygen partial pressure is higher than that at high oxygen partial pressure. The active oxidation of SiC led to a small amount of silica formation, which in turn increased the inward transportation of oxygen, resulting in enhanced oxidation of the material. Similar results were observed by Rezaie et al. [17]. They reported that no protective oxide layer was formed on ZrB2-SiC at an oxygen partial pressure of 10\\x0010 Pa due to the active oxidation of SiC at this oxygen partial pressure. The oxidation resistance of ZrB2-SiC was mostly dependent on the SiC content. A high SiC content is for oxidation resistance at 1800 °C in air due beneficial to a large amount of silica glass formation (Fig. 4). However, this phenomenon changed with reduced oxygen partial pressure. The high SiC content is not responsible for the improved oxidation resistance of the material at the lower oxygen partial pressure since the preferential active oxidation of SiC resulted in little silicate formation and large amounts of gaseous oxidation products (i.e. CO, SiO). The formation of the gaseous phase products, especially at high partial pressures, would accelerate the volatilization of the initially formed glass and create short-circuit paths for the incoming oxygen, resulting in increased oxidation rates. Figure 6 shows the morphologies of the SiC-depleted region, which are parallel and perpendicular to the top surface of the oxidized specimens. Note that SiC forms a network that is interconnected in three dimensions, as shown in Figure 6, although it is discontinuous in two dimensions, as can be seen in Figure 1. Moreover, the formation of the SiC-depleted layer resulted from its interconnectivity in three dimensions since the diffusion of oxygen in ZrB2 is very low. Apparently, the degree of SiC interconnectivity in the matrix increased with increasing SiC content. A high degree of SiC interconnectivity causes the rapid active oxidation of SiC. Therefore, the thickness of the SiC-depleted layer for oxidized ZS2 is higher than that for oxidized ZS1 under the same conditions. In this respect, the degree of SiC interconnectivity in the matrix should be reduced by optimization of the composition and microstructure for improved oxidation resistance. According to the above analysis, a high SiC content is beneficial for the oxidation resistance of the material at high oxygen partial pressure whereas it is detrimental at low oxygen partial pressure. We should therefore make a trade-off between them in response to the practical application conditions, and the degree of the SiC interconnectivity in the matrix should be reduced by optimizing the composition and microstructure. In conclusion, dense ZrB2-SiC UHTCs were prepared by hot-pressing. ZrB2 containing 30 vol.% SiC particulates exhibited higher oxidation resistance at 1800 °C under an oxygen partial pressure of 0.2 atm compared with ZrB2 + 20 vol.% SiC. High SiC content  Figure  6. Morphologies  of  the  SiC-depleted  region, which  are  (a)  parallel and (b) perpendicular to the top surface of the oxidized ZS2.  for oxidation resistance at 1800 °C in air is beneficial due to a large amount of silica glass formation. However, this trend is reversed at a reduced oxygen partial pressure due to rapid active oxidation of SiC resulting in little formation of silica glass. In fact, SiC forms a network that is interconnected in three dimensions, although it is discontinuous in two dimensions. The degree of SiC interconnectivity in the matrix increased with increasing SiC content. In this respect, the degree of the SiC interconnectivity in the matrix should be reduced by optimizing the composition and microstructure for improved oxidation resistance. High SiC content is beneficial for the oxidation resistance of the material at high oxygen partial pressure whereas it is detrimental at low oxygen partial pressure. Therefore, we should make a trade-off between them in response to the practical application conditions.  This work was supported by the National Natural Science Foundation of China (90505015 and 50602010), the Research Fund for the Doctoral Program of Higher Education (20060213031) and the Program for New Century Excellent Talents in University.  [1] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, J. Eur. Cer. Soc. 22 (14-15) (2002) 2757. [2] K. Upadhya, J.M. Yang, W.P. Hoffman, Am. Ceram. Soc. Bull. 76 (12) (1997) 51. [3] F. Monteverde, A. Bellosi, Adv. Eng. Mater. 6 (5) (2004) 331. [4] W.G. Fahrenholtz, G.E. Hilmas, A.L. Chamberlain, J.W. Zimmermann, J. Mater. Sci. 39 (2004) 5951. [5] D.M. Van Wie, D.G. Drewry Jr., D.E. King, C.M. Hudson, J. Mater. Sci. 39 (19) (2004) 5915. [6] D.M. Van Wie, D.J. Risha, C.F. Suchomel, AIAA Paper, 42nd AIAA Aerospace Sciences Meeting and Exhibit (2004) 11101. [7] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, J. Mater. Sci. 39 (19) (2004) 5887. [8] F. Monteverde, A. Bellosi, (2005) 1025. [9] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, Refrac. Appl. Trans. 1 (2) (2005) 1. [10] F. Monteverde, Corros. Sci. 47 (8) (2005) 2020. [11] F. Monteverde, A. Bellosi, Solid State Sci. 7 (5) (2005) 622. [12] S.R. Levine, E.J. Opila, NASA/TM—2003-212483. [13] E.J. Opila, S.R. Levine, J. Lorincz, J. Mater. Sci. 39 (19) (2004) 5969. [14] S.S. Hwang, A.L. Vasiliev, N.P. Padture, Mater. Sci. Eng. A (2007), doi:10.1016/j.msea.2007.03.002. [15] W.C. Butterman, W.R. Foster, Am. Mineral. 52 (1967) 880. [16] Alireza Rezaie, W.G. Fahrenholtz, G.E. Hilmas, J. Eur. Cer. Soc. 27 (2007) 2495. [17] Alireza Rezaie, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Cer. Soc. 89 (10) (2006) 3240.  J. Eur. Ceram. Soc.  25  (7)  \\x0c']"
},{
  "_id": 148,
  "PDF": "Oxidation behavior of ZrB2 composites doped with various transition metal silicides.pdf",
  "Text": "['Corrosion Science 83 (2014) 281-291  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Oxidation behavior of ZrB2 composites doped with various transition metal silicides  ⇑  Laura Silvestroni  , Giacomo Meriggi, Diletta Sciti  CNR-ISTEC,  Institute of Science and Technology for Ceramics, Via Granarolo 64,  I-48018 Faenza,  Italy  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 19 December 2013 Accepted 12 February 2014 Available online 20 February 2014  Keywords:  Ceramic SEM Oxidation  1. Introduction  This paper deals with the oxidation behavior of ZrB2-based composites sintered with different additives, namely ZrSi2, MoSi2, TaSi2 and WSi2. The oxidation mechanisms were investigated between 1200 and 1800 °C for 15 min in a bottom loading furnace. The scope of this study is to draw a classification of goodness for the 4 composites depending on the temperature range and understand how each cation influences the oxidation behavior of ZrB2 by acting either on glass or on ZrO2 modification. MoSi2 was the best additive for improving the oxidation resistance of ZrB2, even up to 1800 °C. Ó 2014 Elsevier Ltd. All rights reserved.  Zirconium diboride, ZrB2, is one of the most investigated material in the ultra-high temperature ceramics (UHTC) class in view of its interesting combination of physico-chemical and engineeristic properties [1,2]. The good mechanical strength of ZrB2-composites has been assessed in various works [3,4] and superior resistance at high temperature has been also reported upon addition of suitable additives [5-7]. However, clear reasons explaining this good behavior at high temperature have not been stated yet. Literature review shows that the introduction of cations of transition metal in form of silicide or carbide, strongly modifies the high temperature performances and the oxidation behavior of ZrB2, but data are not directly comparable owing to different compositions, oxidation conditions and specimen size. The oxidation behavior of monolithic ZrB2 has been object of numerous investigations [8] and all the studies assessed that the main reactions involve the formation of glassy B2O3, at low temperature, and porous crystalline ZrO2, according to (1):  ZrB2 þ 5=2O2 ! B2O3 þ ZrO2  ð1Þ  Boron oxide, tends to evaporate around 1300 °C, depending on the oxygen partial pressure, leaving an open skeleton of ZrO2. This develops in shape of tiny rounded particles on the boride grains and undergoes notable grain coarsening above 1500 °C. It has been stated that the oxidation behavior is dominated by parabolic kinetics despite formation of a porous oxide of the refractory metal [8].  ⇑ Corresponding author. Tel.: +39 546 699723; fax: +39 546 46381. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  http://dx.doi.org/10.1016/j.corsci.2014.02.026 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.  This protective behavior is attributed to the borosilicate glass that fills the pores of the porous oxide scale and often also results in a glassy external layer. The main issues in the oxidation of ZrB2 at temperature below 2000 °C consist indeed in the formation of a porous ZrO2 scale, which allows continuous penetration of oxygen through the bulk, and in the large volume expansion associated to the polytypic transformation of ZrO2 [9]. The addition of SiC [10- 13] or silicides of transition metals [14-16] notably modifies the oxidation behavior of monolithic ZrB2, through the formation of a silico-boride glass, whose viscosity can be modified by the incorporation or dissolution of the relevant cation [14-16], or even through the change of the physical properties of ZrO2, which can melt at lower temperature and form a compact dense external scale [17]. In the present work, several transition metal disilicides, MeSi2, were added to a ZrB2 matrix in order to study how each cation modifies the oxidation performances of this UHTC. The additives object of this study include ZrSi2, MoSi2, TaSi2 and WSi2.  1.1. Effect of MeSi2 on the oxidation of ZrB2 composites  In this paragraph, the effect of transition metal silicides on the oxidation behavior of ZrB2 and ZrB2-SiC composites is briefly reviewed. In a study by Lavrenko et al. [18] it was assessed that the addition of ZrSi2 improved the oxidation behavior at 1200 °C of monolithic ZrB2 and further surpassed the MoSi2 and WSi2 hightemperature materials in oxidation resistance. The effect of MoSi2 on the oxidation resistance of ZrB2 has been studied through thermogravimetric analyses up to 1400 °C for 30 h revealing parabolic kinetics at 1400 °C, due to the formation of a compact silica-based  \\x0c', '282  L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  scale and subsurface oxidation with formation of refractory species, like MoB [19]. Similar oxidation products were observed upon oxidation at 1400 °C for 100 h in conventional kilns [20]. A successive work reported different oxidation stages for HfB2 doped with TaSi2, revealing formation of a layered oxide scale based on Hf6Ta2O17 crystals embedded in a silica layer upon oxidation at 1600 °C for 15 min in conventional air furnace [21]. In a very recent paper [22], the oxidation behavior of ZrB2 with 0, 4, 6, or 8 mol% W was investigated at 800-1600 °C and pointed out that the addition of W into B2O3 increased the stability of the protective glassy layer, which resulted in higher oxidation resistance. Most of the works in the literature, consider the effect of silicides addition in ZrB2-SiC composites, however, being SiC a major phase, in amount 10-40 vol%, it is difficult to discriminate the action of the silicide alone. For example, the effect of ZrSi2 on the oxidation resistance of ZrB2-SiC ceramics has been studied up to 1700 °C in air by thermogravimetric analysis coupled with differential thermal analysis and it has been concluded that addition of small amounts of this silicide, 7 vol%, substantially slow down the oxidation of the ceramic due to the stabilization of the upper oxide layer of borosilicate glass [23]. A model for ZrB2-SiC-ZrSi2 ceramics during long-term oxidation at 1500 °C has been also developed and validated by experiments, confirming a parabolic behavior. The reason ascribed to the improved performances was related to decreased oxygen diffusion rate in the subsurface layer [24]. ZrB2-SiC-MoSi2 ceramics were studied at 1500 °C for 10 h and the results confirmed good oxidation behavior, as previously pointed out for ZrB2-SiC and ZrB2-MoSi2 composites [25]. Concerning the effect of tantalum, other authors [26-29] pointed out that Ta-addition is beneficial up to 1600 °C by acting both on zirconium oxide and glass properties. As for the first effect, Ta, with a higher valence cation than Zr, can stuff oxygen vacancy in Zr oxide. The presence of fewer vacancies in the refractory oxide makes it more stable and thus the tetragonal/monoclinic phase transformation is circumvented together with the volume expansion associated, and thus the possibility to form cracks in the scale is reduced too. As far as the glass is concerned, the borosilicate glass becomes more viscous due to the higher cation strength field of tantalum inducing immiscibility and the formation of a more viscous external layer possessing higher boiling point and preventing oxygen inward diffusion. However tantalum is reported to become detrimental above 1600 °C [28] due to the melting of Ta2O5 and to the formation of a complex orthorhombic Zr-Ta-oxide which has a needle-like morphology not favorable for adhesion to the unoxidized bulk. As for WSi2 additions, there are no reports on its effect on the oxidation. However, beneficial effects of W-compounds on the oxidation behavior of ZrB2 and ZrB2-SiC composites have been already outlined in [6,22,30]. W-doping was introduced in ZrB2-SiC composites in form of carbide in amounts ranging from 3 to 7 vol% and increased performances of the W-doped ceramics as compared to pure ZrB2-SiC was reported in all cases. The reasons ascribed to this improvement rely on the formation of an eutectic between WO3 and ZrO2 at 1275 °C, that acted as liquid phase sintering and decreased the porosity of the scale [6,30]. Another hypothesis is that WO3 itself acted as barrier due to the volume increase associated with oxidation of W to WO3 [30], or W increased the stability of the borate glass retarding its evaporation [22]. Other authors suggested that tungsten oxide species form strong acid sites on ZrO2 and inhibit ZrO2 crystallite tetragonal to monoclinic structural transformations. [31]. The scope of this work is hence a direct comparison of the effect of various silicides on the oxidation of ZrB2. The rationale behind this work is that a better understanding of the basic oxidation mechanisms of ZrB2 doped with MeSi2 will lead to the development of  better and more refractory composites, or to the definition of safety temperature ranges for each composite.  2. Experimental procedure  ZrB2-based composites were prepared using the following commercial powders: hexagonal ZrB2, Grade B (H.C. Starck, Germany), specific surface area 1.0 m2/g, impurity max content: C: 0.25 wt%, O: 2 wt%, N: 0.25 wt%, Fe: 0.1 wt%, Hf: 0.2 wt%, particle size range 0.1-8 lm; orthorhombic ZrSi2-F (Japan New Metals Co., LTD., Osaka, Japan), particle size 2-5 lm, impurities: C > 0.15, Fe > 0.30, O > 1.00; tetragonal MoSi2 (<2 lm, Aldrich, Steinbeim, Germany), particle size range 0.3-5 lm and oxygen content \\x181 wt%; hexagonal TaSi2 (ABCR, GmbH & Co., Karlsruhe, Germany), particle size <45 lm; tetragonal WSi2 (Sigma Aldrich, Milano), \\x00325 mesh, 99.5%, traces of metals <6000 ppm. Compositions containing 15 vol% of silicide were prepared by conventional wet milling route and sintered by hot pressing to full density, as reported in [5]. For the sake of clarity, Table 1 summarizes the sintering conditions and the main microstructural features of the 4 composites. Oxidation tests were carried out in a bottom loading furnace (Nannetti FC18, Faenza, Italy and Carbolite BLF1800 furnace, Hope Valley, UK) at 1200, 1350, 1500 and 1650 °C for 15 min on rectangular 13 mm \\x02 2.5 mm \\x02 2 mm bars in static air. For the best compositions, further tests at 1800 °C were performed. Specimens were located in the furnace when the maximum temperature was achieved and then removed and air-quenched after the exposure time. The mass and dimensions of the bars were measured before and after the oxidation. The as-sintered material and oxidized specimens were examined on the surface using X-ray diffraction (Bruker D8 Advance, Bruker, Karlsruhe, Germany) to identify the crystalline phases. The microstructures before and after oxidation were analyzed using field emission gun scanning electron microscopy (FE-SEM, Carl Zeiss Sigma NTS Gmbh Öberkochen, Germany) and energy dispersive spectroscopy (EDS, INCA Energy 300, Oxford Instruments, UK) on surface and fractured cross-section of the specimens to reveal modifications induced by oxidation.  3. Results  3.1. As sintered microstructure  In order to understand the microstructural evolution upon oxidation at the various temperatures, it is useful to briefly summarize the compositional structure of the as sintered composites. For some materials, a thorough characterization can be found in [5,32], but the main features are reported in Table 1. ZBZ - This material was fully densified at 1600 °C. The polished surface of ZBZ is reported in Fig. 1a, evidencing mean grain size of 2.5 lm and the secondary phases (vol%): 3.2 ZrO2, 3.0 ZrSi2, 1.5 SiO2-based species and 1.3 SiC. No wetted grain boundaries were noticed by SEM and TEM inspections. Often, SiO2 pockets surrounding elongated SiC crystals were found. ZBM - This ceramic achieved the full density at 1750 °C, Table 1. In Fig. 1b, ZrB2 had mean grain size around 2.5 lm and the matrix had a peculiar morphology with ZrB2 grains displaying a core-shell substructure. The core was constituted by original ZrB2 grains and the shell by a (Zr,Mo)B2 solid solution, with an amount of Mo around 5 at.%, which grew epitaxially on the core [32]. The phase with white contrast, often located at the tips of MoSi2, is MoB or Mo-Si-B phase, in amount below 2%. Silica pockets, around 4 vol%, were recognizable as dark contrasting phases. Small amounts, below 2%, of ZrO2, ZrC, and spurious Zr-C-O phase were also detected.  \\x0c', 'L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  283  Table 1 Composition, sintering conditions and main microstructural  features of the as sintered ZrB2-MeSi2 composites. m.g.s.: mean grain size.  Label  ZBZ ZBM ZBT ZBW  MeSi2 15 (vol%)  Sintering (°C, MPa, min)  Rel. density (%)  ZrB2 m.g.s. (lm)  Secondary phases (vol%)  ZrSi2 MoSi2 TaSi2 WSi2  1600, 30, 10 1750, 30, 20 1850, 30, 10 1930, 30, 25  96.5 99.0 99.0 95.5  2.5 2.4 2.0 3.5  3.2 ZrO2, 3 ZrxSiy, 1.5 SiO2, 1.3 SiC (Zr,Mo)B2, 14.2 MoSi2, 4.0 SiO2, 1.5 SiC, 1.0 ZrO2 (Zr,Ta)B2, 5.0 TaxSiy, 2.0 SiO2, 1.6 ZrO2, 1.5 SiC (Zr,W)B2, 3.5 [WC, (Zr,W)B, (W,Zr)xSiy], 2.4 SiO2, 0.5 ZrO2  Fig. 1. Microstructure of the as sintered composites. (a) ZBZ, (b) ZBM, (c) ZBT, (d) ZBW.  ZBT - This composite required 1850 °C to achieve the full density. An example of polished surface of ZBT is shown in Fig. 1c. ZrB2 grains had an average size of about 2 lm and, similarly to ZBM, these grains were surrounded by a (Zr, Ta)B2 solid solution, which appeared lighter in color. According to quantitative EDS analysis and XRD, the composition of this solid solution was (Zr0.8Ta0.2)B2 [5]. Pure TaSi2 was not clearly identified in the composite, but bright TaxSiy phases were observed in amount around 5 vol%. About 1.6 vol% of ZrO2 particles were also found along with 2 vol% of silica-based phases containing various impurities. Spurious carbide phases, such as Zr-Ta-C, SiC, or Si-C-O, were also detected in limited amounts, below 2 vol%. ZBW - 1900 °C were necessary to densify this composite, owing to a more refractory silicide. The microstructure of the material sintered with WSi2 is depicted in Fig. 1d. Also this composite displayed a core-shell morphology: pure ZrB2 constitutes the core and a (Zr, W)B2 solid solution the shell, with a W content around 2 at.% and mean grain size in the order of 3.5 lm. Large agglomerates of white phases, in amount around 3.5 vol%, were identified as WC and W-C-O with (Zr, W)xSiy and (Zr, W)B as terminal parts in triangular shape. About 2.8 vol% of SiO2 pockets trapping ZrO2 particles were observed too.  3.2. Microstructure upon oxidation  3.2.1. Appearance and weight gain  A picture of the specimen after oxidation tests between 1200 and 1800 °C is shown in Fig. 2, where it can be seen that increasing the oxidation temperature, the grey color, typical of the as sintered bars first on the left, moves to dark grey, index of glass formation, to whitish-yellow, indicating the presence of mostly ZrO2. It is worthy to note that ZBM was the least damaged material, preserving  the darkest color even at 1800 °C. On the contrary, the most damaged sample was ZBT, which finished the 1650 °C test with compromised shape and evident bubbling on the surface. The weight gain per unit surface area as a function of the oxidation temperature, plotted in Fig. 3, increased with the temperature for all composites and confirmed the results suggested by visual inspection: even after oxidation at 1800 °C, ZBM was the least damaged and ZBT the most damaged at 1650 °C. Up to 1500 °C the material gaining the highest weight was ZBZ, followed by ZBW, ZBT and ZBM. Above this temperature the trend changed: ZBT abruptly increased, ZBZ and ZBW were almost comparable, whereas ZBM arrived to the highest temperature with a weight gain around one third as compared to the other composites.  3.2.2. X-ray diffraction  Table 2 reports the crystalline phases detected by XRD on the various specimens at the different oxidation temperature. The main crystalline oxidation product in ZBZ was monoclinic zirconia at all temperatures, the tetragonal polytype was present too but in minor amount and ZrB2 peaks remained evident up to 1350 °C. Also for ZBM composite, the main crystalline phase was monoclinic ZrO2, but ZrB2 reflections remained visible even after oxidation at all temperatures, although just in traces above 1650 °C. MoSi2 peaks disappeared after oxidation at 1350 °C leaving instead tetragonal MoB phase. The high temperature ZrO2 polytype appeared after oxidation at 1350 °C and increased in intensity with the oxidation temperature. No clear Mo5Si3 phase was identified. At 1800 °C, besides monoclinic and tetragonal ZrO2, MoO3 was also clearly identified. The X-ray diffraction patterns for ZBT were quite diverse depending on the oxidation temperature, as Fig. 4 shows. At 1200 °C, monoclinic ZrO2 was the major phase, but peaks relative to the boride  \\x0c', '284  L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  Fig. 2. From left to right appearance of the oxidized samples at 1200, 1350, 1500 and 1650 °C for 15 min. (a) ZBZ, (b) ZBM, (c) ZBT, (d) ZBW. For ZBM and ZBW also the specimens oxidized at 1800 °C are reported in the bottom.  ZBZ - The external surface of ZrB2 doped with ZrSi2 after oxidation at 1200 °C was completely covered by a continuous borosilicate glass from which 500 nm ZrO2 grains emerged from time to time, as Fig. 5a shows. An increase of the oxidation temperature to 1350 °C resulted in a thickening of the external silica glass, where boron was not detected anymore and the ZrO2 grains were less abundant, Fig. 5b. The same phases were observed after test at 1500 °C, but the amount of protruding ZrO2 increased and assumed the shape of platelets, petals or tiny rods, Fig. 5c. At 1650 °C the surface was very similar to that observed at 1500 °C, but in the edges region the petal-like ZrO2 was more frequent than the horns, which was instead more abundant in the center of the specimen, Fig. 5d. Images of the cross-section of ZBZ are reported below in Fig. 5. Overall, the oxidized sections reveal compact scales with no porosity. Already at 1200 °C, Fig. 5a, about 1 lm thick of continuous silica layer covered a 18 lm sublayer of rounded fine zirconia. In the following 10 lm inward, bulk material with intergranular zirconia around ZrB2 grains was found. At 1350 °C, Fig. 5b, the external silica layer thickened just to 2 lm, then a 15 lm of coarse ZrO2 filled with silica was standing above 30 lm of ZrB2 with intergranular fine ZrO2. Increasing the oxidation temperature to 1500 °C, Fig. 5c, the outermost 10 lm were composed by SiO2 glassy phase topping ZrO2 crystals. Right beneath this scale, 100 lm of dense ZrO2 phase with SiO2 pockets was found. The last oxidation temperature, 1650 °C, Fig. 5d, resulted in the formation of a 7 lm dense silica scale from which ZrO2 rods emerged on the top (inset in Fig. 5d). Then about 50 lm were constituted by elongated ZrO2 grains filled with SiO2. Moving further inward, the microstructure was mainly columnar ZrO2 for about 120 lm, which became rounded for 30 lm and this phase stood on a mixed ZrB2 scale, around 35 lm thick, containing intergranular ZrO2 and notable amount of SiO2. ZBM - The external surface of ZrB2 doped with MoSi2 after oxidation is shown in upper part of Fig. 6. After oxidation at 1200 °C, a continuous silica layer formed on the top and tiny crystal of ZrO2 could be occasionally found immersed in, Fig. 6a. Traces of molybdenum oxide and evidence of Ca-rich immiscible glass could be found as well. Oxidation at 1350 °C induced the formation of higher amount of ZrO2 precipitates on the continuous silica scale,  Fig. 3. Weight gain per unit surface area as a function of the oxidation temperature.  were well visible and those corresponding to the (Zr, Ta)B2 solid solution seemed more intense as compared to the pure ZrB2, in addition TaB2 peaks and mixed Zr-Ta oxide with orthorhombic structure and formula TaZr2.75O8 were found too. At 1350 °C, the main phase was still m-ZrO2, but TaB2 peaks clearly emerged whilst ZrB2 reflections decreased in intensity. When the composite was oxidized at 1500 °C, the mixed oxide was the main phase, which showed evident crystal orientation along the [2 0 0] planes, monoclinic ZrO2 and TaB2 followed. Increasing the oxidation temperature to 1650 °C the preferential orientation of TaZr2.75O8 changed along the [0 2 0] planes and, besides ZrO2, orthorhombic Ta2O5 formed. Moving to ZBW composite, the main crystalline phase was monoclinic ZrO2, in addition tungsten oxide was detected at all temperatures. Traces of WB were instead visible only up to 1350 °C, whilst ZrB2 reflections were not identified at any oxidation temperature.  3.2.3. SEM-EDS analyses  In the following, the microstructural details observed on the external surface and on the fractured cross-section will be presented for each composite, Figs. 5-8. Scale thickness as calculated from SEM analyses is reported in Fig. 9.  Table 2 Main crystalline phases detected by XRD on the oxidized specimens. Compounds are listed in decreasing amount.  Sample  1200 °C  1350 °C  1500 °C  1650 °C  m-ZrO2, ZrB2, t-ZrO2 m-ZrO2, ZrB2, MoB, MoSi2  m-ZrO2, t-ZrO2, ZrB2 m-ZrO2, ZrB2, MoB, t-ZrO2  a  m-ZrO2, t-ZrO2 m-ZrO2, ZrB2, MoB, t-ZrO2  m-ZrO2, t-ZrO2 m-ZrO2, MoB, t-ZrO2, ZrB2  a  m-ZrO2, ZrB2, o-Zr2.75TaO8, TaB2 m-ZrO2, c-WO3, t-WB  m-ZrO2, TaB2, o-Zr2.75TaO8, ZrB2 m-ZrO2, c-WO3, t-WBa  o-Zr2.75TaO8 [2 0 0], m-ZrO2, TaB2 m-ZrO2, c-WO3  o-Zr2.75TaO8 [0 2 0], o-Ta2O5, mZrO2 m-ZrO2, c-WO3  ZBZ ZBM  ZBT  ZBW  a Traces.  1800 °C  - m-ZrO2, t-ZrO2, MoO3 -  m-ZrO2, c-WO3  \\x0c', 'L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  285  Fig. 4. X-ray diffraction patterns of ZBT after oxidation at each temperature for 15 min.  Fig. 5. SEM images of surface (up) and cross-section (bottom) of ZBZ oxidized at (a) 1200, (b) 1350, (c) 1500, (d) 1650 °C with microstructural details in the insets.  Fig. 6b. The same microstructural features were observed after oxidation at 1500 °C, where preferential glass removal at the edges was observed, Fig. 6c. The same holds true for oxidation at 1650 °C, Fig. 6d, where ZrO2 nanoprecipitates and aggregates were floating on continuous silica scale. The surface of the specimen oxidized at 1800 °C still showed a continuous glassy scale with some very large transparent bubbles, Fig. 2b. SEM inspection revealed the silica glass in the dark areas and zirconia where the bubble exploded. At the glass-zirconia interface, MoO3 grains in platelet shape were observed, as illustrated in Fig. 6e. Images of the cross-section of ZBM are reported in the bottom part of Fig. 6. After oxidation at 1200 °C, Fig. 6a, a continuous silica layer was covering about 2.5 lm of fine graded ZrO2 and right underneath, MoSi2 oxidation products, like SiO2 pockets coasted by tiny MoB/Mo5Si3 particles, were found in ZrB2. A temperature increase of 150 °C left an almost unaffected thickness of silica and zirconia, but induced the migration of further silica close to the upper surface in a spin shape, Fig. 6b, indicating that at this temperature silica is very fluid. The cross-section of ZBM at 1500 °C in Fig. 6c shows a thickening of ZrO2 layer, whilst the continuous external coating remained almost around a couple of  micron thick. Underneath the silica spin, SiO2 and MoB could be easily found in a striped shape always adjacent one another immersed in ZrB2 and the total altered thickness passed from 4 lm at 1200 °C to 25 lm at 1500 °C. At 1650 °C, Fig. 6d, the outermost silica layer was crossed through by zirconia grains and its thickness increased to about 25 lm, the ZrB2-SiO2-MoB layer resulted around 40 lm and then the bulk ceramic started. The cross-section of ZBM oxidized at 1800 °C is shown in Fig. 6f. The outermost scale composed by 50 lm of is compact ZrO2 and silica with some cavities due to bubbling. Underneath this layer, bright MoB chain separates other 65 lm of columnar ZrO2 crossed by silico-boride glass. Then ZrB2 starts with SiO2 and MoB phases, similarly to the previous samples before finding the unreacted bulk. ZBT - The external surface of this ceramic after oxidation in the 1200-1650 °C temperature range is shown in Fig. 7, upper part. At 1200 °C, the specimen was covered by a continuous SiO2-based layer with B2O3 traces and tiny aggregates of ZrO2 particles. No Ta was detected both in the glass and in the crystalline particles. At 1350 °C the ZrO2 aggregates increased in dimensions from 500 nm to 6 lm and boron was not detected anymore in the glass. The surface after exposure at 1500 °C appeared completely  \\x0c', '286  L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  Fig. 6. SEM images of surface (up) and cross-section (bottom) of ZBM oxidized at (a) 1200, (b) 1350, (c) 1500, (d) 1650 °C with microstructural details in the insets. (e) External surface and (f) cross-section of ZBM oxidized at 1800 °C with microstructural details as indicated.  Fig. 7. SEM images of surface (up) and cross-section (bottom) of ZBT oxidized at (a) 1200, (b) 1350, (c) 1500, (d) 1650 °C with microstructural details in the insets.  \\x0c', 'L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  287  Fig. 8. SEM images of surface (up) and cross-section (bottom) of ZBW oxidized at (a) 1200, (b) 1350, (c) 1500, (d) 1650 °C with microstructural details in the insets. (e) Surface and (f)-(g) cross-section of ZBW oxidized at 1800 °C with microstructural details as indicated.  The cross-sections of the oxidized ZBT samples are reported in Fig. 7, bottom part. The sample oxidized at 1200 °C involved less than 10 lm thick with an outermost continuous silica scale, a thin layer of fine grained ZrO2 and other about 5 lm of boride matrix containing the oxidation products of the sintering additive, like SiO2 and Ta2O5, Fig. 7a. At 1350 °C the silica layer slightly thickened and tended to detach from the underlying zirconia, which involved around 20 lm of material, Fig. 7b. In this scale, SiO2 and Ta2O5 were still present and recognizable as dark and white phases, respectively. TaB2 was present too as white phase. At 1500 °C, Fig. 7c, the silica scale thickness doubled and completely surrounded the Zr-Ta-oxide crystals; differently from the previous sample, no spalling was observed between glass and zirconia. Underneath this layer, about 25 lm of ZrO2 with bright tantalum oxide and boride white agglomerates were found next to silica pockets. The oxidation at 1650 °C provoked catastrophic microstructural evolution in ZBT, Fig. 7d. About 500 lm from the external surface were composed by ZrO2 and Ta2O5 partially filled with silica glass where bubbles and craters involved the whole thickness. The ZrO2 grains were coated with dendritic Ta2O5, as shown in the upper inset of Fig. 7d, and in the region right underneath these crystals grew in a spike-pillar shape, like reported in the second inset of Fig. 7d. This thick scale was standing above a 10 lm layer of dense zirconia interposed by bright TaB2 particles standing adjacent to SiO2. Then other 10 lm of material were composed by ZrO2 with intergranular Ta2O5.  Fig. 9. Plot of the modified thickness in the ZrB2-based samples as a function of the oxidation temperature.  changed: most of the surface was occupied by a crystalline mixed oxide, ZrTa2.75O8, with petal-like shape, and this structure was filled by silica. The crystals were in turn covered with brighter dendritic structures richer in tantalum. At 1650 °C the sample underwent a second drastic change: the mixed Zr-Ta-oxide was lifted by the outgassing volatile species and the platelets assumed a vertical position immersed in the silica-based glass. Large craters were present too. This mixed oxide was disposed in geometrical shape and was covered by small dendritic Ta2O5 crystals, Fig. 7d.  \\x0c', '288  L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  ZBW - The external surfaces of the sample containing WSi2 oxidized at all temperatures are reported in Fig. 8, upper part. The surface of the specimen oxidized at 1200 °C appeared covered by B2O3 and silica glass where sub-micrometric zirconia grains were dispersed in. These were covered by a W-O-based dendritic network and interposed with 500 nm aggregates of bright WO3 grains, Fig. 8a. At 1350 °C boron oxide disappeared and the surface remained composed by ZrO2 grains immersed in silica glass. Also the W-signal was barely detected among the crystals. The oxidation at 1500 °C induced a notable variation of the external morphology which appeared wavy and rutty, showing bumps and craters, Fig. 8b. The ZrO2 grains coarsened to about 2 lm and appeared aggregated and melted on a silica continuous layer. The growing planes were well visible in ZrO2 grains and W-rich phases with bright contrast were found at the grain boundary, Fig. 8c. The aspect of the sample oxidized at 1650 °C was not notably different from the one oxidized at 1500 °C, just more bumps and craters were noticed, Fig. 8d. An increase of the oxidation temperature up to 1800 °C provoked notable damage to the ZBW sample, it was completely covered by white oxide and SEM analyses evidenced cracking and vigorous bubbling in the ZrO2-SiO2 outermost scale, Fig. 8e. The cross-sections of the oxidized ZBW samples are shown in Fig. 8, bottom part. After oxidation at 1200 °C a discontinuous silica scale was standing on top of a fine-grained zirconia scale where bright rounded W-O-based particles were dispersed in, Fig. 8a. This scale, less than 25 lm thick, was well adherent to the unreacted bulk. An increase of the oxidation to 1350 °C doubled the thickness of the oxidized scale as compared to the previous sample, which maintained however the same morphology, Fig. 8b. Also in this case the oxidized layer was well anchored to the matrix, thanks to the growth of the oxide well integrated onto the boride grains, inset in Fig. 8b. The oxidation at 1500 °C, Fig. 8c, induced a diversification in morphologies of the ZrO2 sub layer. Starting from the top, 20 lm coarse ZrO2 grains were standing on other 20 lm columnar ZrO2, which was topping about 95 lm of dense rounded-grains ZrO2. The whole thickness was partially filled with silica and the oxide grains were coated with homogeneously dispersed WO3 drops which eventually turned into agglomerates where enough porosity among ZrO2 grains let enough room for their condensation as 5 lm particles, see the insets in Fig. 8c. At 1650 °C the morphology of the various scale remained basically unmodified as compared to the features observed for the samples oxidized at 1500 °C, Fig. 8d, but each scale just increased in thickness and columnar ZrO2 layer further developed. Another feature worthy of mentioning is the formation of small cracks within columnar ZrO2 grains. At 1800 °C the sample resulted much damaged, about 400 lm of material underwent oxidation, Fig. 8f. The outermost scale was composed by large ZrO2 grains fissured and partially filled with silica-based glass. This thick layer was standing on top of a denser scale made of ZrO2 with elongated shape filled by white WO3 particles and veined by B2O3 glass, Fig. 8g. The presence of borates at such high temperature could be due to the stabilization of the glass by W, as reported in [22]. This scale was completely spalled from the bulk.  4. Discussion  An impression at glance of the beneficial or detrimental effect of each silicide is provided in Fig. 9 which reports the modified thickness of the four composites after oxidation, where for oxidized thickness we intend the depth of material containing oxidation products, either of the diboride, or the silicide. The plot is in good agreement with the weight gain in Fig. 3, i.e. larger weight gain is associated to a thicker oxide layer, suggesting that any of the two  indicators, weight gain or thickness, could be used as representative of oxidation degree. Multiple and various phenomena occur upon oxidation of ZrB2 doped with various MeSi2 implying melting, formation of glassy phases, phase transformation and volatilization. Such mechanisms depend on the chemico-physical nature and content of the silicide after sintering. The oxidation products of the silicides are also function of the valence of the cation, which will form solid, volatile or instable oxides and will interact with the boride matrix and its oxidation products. Figs. 3 and 9 and microstructural characterization evidence similarities, but also strong differences amongst the systems analyzed. All the systems develop a complex multilayered scale basically containing ZrO2-based oxides, whilst the silicides provide some partially protective silica-based oxide scales. Strong oxidation phenomena involve the three systems from 1650 °C on, whilst ZBM remains more unaffected from oxidation. In the subsurface layer, no depletion layer is ever observed for any of these systems. Remarkably, addition of MoSi2 and TaSi2 retard the formation of columnar ZrO2, which is typical of ZrB2 or ZrB2-SiC oxidation [8,10-13]. Columnar ZrO2 is observed for ZBZ and ZBW starting from 1500 °C, whilst for ZBM appears only upon oxidation at 1800 °C. In the following, discussion will compare the materials response at each temperature, keeping in mind three fundamental aspects affecting the oxidation processes:  a. Content, composition and chemical stability of the silicide phase after sintering, as silica-forming phase. Due to the silicide reactivity during sintering the chemico-physical nature and content of silicide after densification can be significantly different compared to the starting composition, as illustrated in paragraph 3.1. Low melting silicides are for instance present in ZBZ [33]. In ZBZ, ZBT and ZBW the final content of silicide is drastically reduced from 15 to 3- 5 vol%, whilst in ZBM remains basically unchanged, 14 vol%, indicating that MoSi2 is the most stable silicide with the lowest tendency to dissociate. b. The type of cation introduced with the silicide. Possible cations effects on glass modification have been already outlined in the paragraph 1.1. In this work, Mo, Ta or W are potentially able to modify the glass composition and viscosity. c. Events occurring in the subsurface layer. Although the surface morphology presents similarities, at least for three composites, the cross-sections reveal different species formed, as the aforementioned columnar ZrO2.  the four systems reveals the following  A direct comparison of features. 1200 °C - At this temperature, all the four composites display similar behavior, with an outer borosilicate layer where ZrO2 particles are dispersed in and this layer is generally continuous. The four silicides are expected to oxidize according to:  ZrSi2 þ 3O2 ! ZrO2 þ 2SiO2  5MoSi2 þ 7=2O2 ! Mo5 Si3 þ 7SiO2  2TaSi2 þ 13=2O2 ðgÞ ! Ta2O5 þ 4SiO2  WSi2 þ 7=2O2 ðgÞ ! WO3 þ 2SiO2  ð2Þ  ð3Þ  ð4Þ  ð5Þ  In this temperature range, a subsurface oxidation layer is formed in all system, mainly composed of ZrO2, for a thickness below 10 lm for ZBM and ZBT, and around 30 lm for ZBZ and ZBW. The different thickness could be related to the actual silicide amount, in the case of ZBM, and to the formation of a more viscous glass in the case of  \\x0c', 'L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  289  ZBT. ZBW seems the most damaged probably owing to the lower content of Si-species. 1350 °C - The picture remains essentially the same, as the four materials still display a similar behavior even if ZBZ and ZBW exhibit a thicker oxide scale compared to the other systems. This is due to the starting of formation of a subsurface layer in ZBM and ZBT containing solid and stable products at these temperatures, like MoB and TaB2, as discussed below. 1500 °C - The four materials begin to remarkably differentiate their oxidation behavior. ZBZ and ZBW start to oxidize readily and both weight gain and oxide thickness reach similar values. ZBM and ZBT have lower oxidation rates, ZBM being the most resistant. This behavior is probably due to the refractoriness of secondary phases. The composite doped with ZrSi2 contains an amount of non-stoichiometric zirconium silicides with low melting points, around or even lower than 1500 °C. Thus in the bulk, liquid phases can arise from ZrxSiy melting. For WSi2 additions, liquid phase can be already present at temperatures lower than 1300 °C, due to the aforementioned WO3- ZrO2 eutectic [34]. Additional sources of liquid come from mixed (Zr, W)-silicides, oxides and carbides present as secondary phases in the sintered microstructure [35]. Liquid phases, both in ZBZ and ZBW bulk, provide a fast path for oxygen transport and enhance oxidation. For ZBW composite, the condensation of tungsten oxide from vapor phase is well evident, as this phase is found as nanosized rounded particles on ZrO2 grains, Fig. 8c and d, therefore tungsten oxide is supposed to come from reactions (5), (6). As a matter of fact, differently from ZBT, the formation of solid solutions in ZBW between W and zirconia was not proved by SEM-EDS nor X-ray diffraction, so reaction (6) is possible, too.  ðZr; WÞB2 þ O2 ! ZrO2 þ WO3  ð6Þ  The addition of MoSi2 has been previously discussed [19,20]. First of all, secondary phases such as MoB and Mo5Si3 are already present in small amounts after sintering, but, differently from residual ZrxSiy detected in ZBZ, they and are highly refractory with melting point overpassing 2000 °C [36]. Oxygen penetration in the subsurface layer favors further formation of these species, according to reactions (3) and (7), where B2O3 in (7) derives from ZrB2 oxidation in (1).  2MoSi2 þ B2O3 þ 5=2O2 ! 4SiO2 þ 2MoB  ð7Þ  The addition of TaSi2 at this temperature showed beneficial effects on the oxidation behavior, limiting the weight gain and also the degradation of a thick scale. This is probably due to similar mechanisms taking place for MoSi2. Indeed, TaB2 and SiO2 were found next to each other’s. Thus, formation of TaB2 can occur according to:  TaSi2 þ B2O3 þ 1=2O2 ! TaB2 þ 2SiO2  ð8Þ  Presence of solid species instead of liquid phases retards the subsurface oxidation, explaining the lower scale thickness of both ZBM and ZBT. In particular, TaB2 is reported to be an excellent oxidation resistance material up to this temperature [37]. TaSi2 is a very peculiar additive, leading to a complex microstructure already upon sintering, thus signing its marked tendency to dissociate and enter in solid solution into ZrB2 lattice. During oxidation it forms silica and an orthorhombic metal oxide according to reaction (4). For this composite, in addition to reaction (1) and (4), also the oxidation reaction of the solid solution has to be considered and can be expressed as (9), bringing to the formation of a mixed orthorhombic oxide:  ðZr; TaÞB2 þ O2 ! TaZr2:75O8  ð9Þ  The chemico-physical properties of this mixed phase are not completely known, however it has been reported to possess lower melting point as compared to that of ZrO2 [28]. The exact mechanism  leading to the formation of this platelet oxide on the surface is not completely understood yet and could also be due to the dissolution of the simple oxides into the silico-boride glass and to the following precipitation as complex oxide, as (10) shows:  11=4ZrO2 þ 1=2Ta2O5 ! TaZr2:75O8  ð10Þ  1650 °C - at this temperature, the aforementioned phenomena are further enhanced. Liquid phases in ZBZ and ZBW cause convection phenomena which make zirconia grow in the form of columns, as consequence of convective transport of low viscosity borosilicate liquid, similar to what observed for ZrB2-SiC composites [10-13]. The outer silica layer, where Zr is almost no soluble in, tends to evaporate and pillar-like ZrO2 features develops on the surface, Fig. 5d, Fig. 8d, owing to the vigorous gas escape. Although WSi2 and ZrSi2 are characterized by different melting points, which should affect the oxidation behavior in different ways, the practical effect of the two silicides up to 1650 °C is very similar in terms of weight gain, scale thickness and oxide morphology. The minimum amount of W that is reported to be effective in promoting liquid phase sintering of ZrO2 and thus hindering the oxidation, was reported to be 4 mol% [17]. In the current case, about 10 mol% are present in the composite so well above the minimum required for densification, but in the investigated temperature range, the superiority of ZBW is not so marked. As a matter of fact, the oxidation protection imparted by W addition is reported to be better above 1600 °C [17] or even above 1800 °C [38]. To the highest step of the podium stands ZBM, which results in the composite with best oxidation resistance at all temperatures. For this material, it seems that the evolution of gaseous phases is prevented up to 1650 °C. Possible reasons could be related to the reduction of B2O3 content in the subsurface liquid through reaction with MoSi2 to form solid MoB, reaction (7). Compared to pure SiO2, the addition of B2O3 decreases the liquid viscosity and its diffusion towards the surface is slower. In addition, the dissolution of little amount of Mo into the silica scale, which is actually reported to acts as a reticulating agent in glass structure [39], could further increase the viscosity of the liquid phases. This combination of mechanisms could explain why formation of columnar ZrO2 is suppressed in this composite. Finally, in ZBT, phenomena similar to those in ZBM could occur, but this is not the case. First of all, the formation of TaB2 did not ensure any further oxidation stability, as this compound was reported to quickly vaporize at 1650 °C [37]. Moreover, the oxidation behavior is dominated by the formation of the mixed Zr, Ta oxide whose formation is accompanied by a notable volume expansion. Another evident feature is the marked gas evolution above 1500 °C which causes the rotation of the TaZr2.75O8 platelets in vertical position, thus keeping open channels to further oxygen inward. The volume expansion, as well as the lower amount of available silica, explains the absence of continuous silica coverage, different from what observed for lower temperature oxidation. Accelerated oxidation led to a marked increase of the weight gain and thickness of the oxidized layer making this as the least resistant composite. The main mechanisms occurring in the four composites in the 1200-1650 °C temperature range are sketched in Fig. 10. 1800 °C - The oxidation in this more severe temperature regime evidenced that in ZBW no stable dense and compact external layer formed, but, on the contrary, continuous migration of silica species and W-oxides to the top surface induced the formation of craters and fissures, owing to a too dynamic gas escape. As for the system containing MoSi2, this temperature eventually induced the formation of columnar ZrO2, owing to boiling of glasses. However the continuous silica scale on the top prevented an extreme weight loss, which remained in the same range or even below that of the other system oxidized at 1500 °C, and, accordingly,  \\x0c', '290  L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  Fig. 10. Sketch of the oxidation phenomena occurring in the ZrB2-based composites during oxidation in the 1200-1650 °C temperature range. (a) ZBZ, (b) ZBM, (c) ZBT, (d) ZBW.  limited the thickness of the modified layer, which did not overpassed that of the other systems oxidized at 1650 °C. The continuous subtraction of B2O3 to preferentially form MoB, instead of diffusing into the silica glass and thus decreasing its viscosity, could be one of the main reasons of the good oxidation behavior of this compound. In addition, the high stability of MoB phase is confirmed by its presence in form of cordillera at the interface between columnar ZrO2 and outer silica scale, first inset in Fig. 6f. It seems that at 1800 °C and at low partial pressure, MoB starts to oxidize according to reaction (11), as we observe its dissolution in silica in the second inset of Fig. 6f.  2MoB þ 9=2O2 ! 2MoO3 þ B2O3  ð11Þ  Then the glass becomes more liquid and induces the formation of ZrO2 pillar. But as long as the oxygen partial pressure returns high enough, MoB phase re-condenses again leaving a more viscous Bdepleted silica glass on the top. Little amount of Mo-based phases managed to stream to the surface and few MoO3 crystals were found, Fig. 6e.  5. Conclusions  ZrB2 ceramics sintered with 15 vol% of various transition metal disilicides, such as Zr, Mo, Ta and W, were oxidized in a bottom loading furnace in the temperature range 1200-1800 °C for 15 min in order to define which additive is better in each temperature regime. The measure of the thickness of the oxidized crosssection and the specific weight gain gave compatible results and the following conclusions were pointed out:  1. Zr, Mo, Ta and W do not clearly modify the outermost glass features, as found by other authors in the case of Ta. 2. WSi2 and ZrSi2 have similar influence on the oxidation behavior of ZrB2, at least up to 1650 °C. The expected superior oxidation resistance in case of WSi2 addition due to ZrO2 scale sintering was not observed even at 1800 °C. 3. TaSi2 could be adequately used up to 1500 °C, but above the abrupt gas evolution and the remarked volume expansion of the mixed oxide render this additive not suitable for ZrB2.  \\x0c', 'L. Silvestroni et al. / Corrosion Science 83 (2014) 281-291  291  4. MoSi2 was by far the best additive to improve the oxidation resistance of ZrB2, due its superior stability during the sintering stage and high temperature exposition in air and due to its ability to form stable solid compounds, like MoB, in the subsurface layer. Limited gaseous species developed and, as a consequence, no formation of columnar ZrO2 was observed up to 1650 °C. The total oxidized scale at 1800 °C was thinner than in the other composites at 1650 °C.  Acknowledgements  D. Dalle Fabbriche (CNR-ISTEC, Faenza), J. Watts and E. Newman (MS&T, Rolla) are greatly acknowledged for oxidation tests. Experiments at 1800 °C have been carried out through the NSF grant ‘‘Materials World Network Program, Cooperative activity in materials research between USA and Italy: Dual composite ceramics for improved properties’’.  References  [6]  [1] R.A. Cutler, Engineering properties of borides, in: S.J. Schneider (Ed.), Ceramics and Glasses: Engineered Materials Handbook, vol. 4, ASM International, Materials Park, 1991, pp. 787-803. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [3] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, High strength ZrB2-based ceramics, J. Am. Ceram. Soc. 87 (2004) 1170-1172. [4] William G. Fahrenholtz, Eric W. Neuman, Harlan J. Brown-Shaklee, Gregory E. Hilmas, Superhard boride-carbide particulate composites, J. Am. Ceram. Soc. 93 (2010) 3580-3583. [5] D. Sciti, L. Silvestroni, C. Melandri, G. Celotti, S. Guicciardi, Sintering and mechanical properties of ZrB2-TaSi2 and HfB2-TaSi2 ceramic composites, J. Am. Ceram. Soc. 91 (2008) 3285-3291. J. Zou, G.J. Zhang, C.F. Hu, T. Nishimura, Y. Sakka, J. Vleugels, O. Van der Biest, Strong ZrB2-SiC-WC ceramics at 1600 °C, J. Am. Ceram. Soc. 95 (2012) 874- 878. [7] Eric W. Neuman, Gregory E. Hilmas, William G. Fahrenholtz, Strength of zirconium diboride to 2300 °C, J. Am. Ceram. Soc. 96 (2013) 47-50. [8] T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, A model for the oxidation of ZrB2, HfB2 and TiB2, Acta Mater. 55 (2007) 5999-6010. [9] H. Okamoto, O-Zr (Oxygen-Zirconium), J. Phase Equil. Diff. 28 (2007) 498. [10] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride - silicon carbide in air up to 1500 °C, J. Eur. Ceram. Soc. 27 (2007) 2495-2501. [11] X.H. Zhang, P. Hu, J.C. Han, Structure evolution oxidation in air, J. Mater. Res. 23 (2008) 1961-1972. [12] P. Hu, W. Guolin, Z. Wang, Oxidation mechanism and resistance of ZrB2-SiC composites, Corros. Sci. 51 (2009) 2724-2732. [13] S.N. Karlsdottir, J.W. Halloran, Oxidation of ZrB2-SiC: influence of SiC content on solid and liquid oxide phase formation, J. Am. Ceram. Soc. 92 (2009) 481-486. [14] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra-high temperature ceramics for aeropropulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, High temperature chemistry and oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3 and TaSi2, J. Am. Ceram. Soc. 91 (2008) 2250-2257.  ZrB2-SiC  [15]  of  during  the  [16] P. Hu, X.H. Zhang, J.C. Han, X.G. Luo, S.Y. Du, Effect of various additives on the oxidation behavior of ZrB2-based ultra-high-temperature ceramics at 1800 °C, J. Am. Ceram. Soc. 93 (2010) 345-349. [17] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Improved oxidation resistance of zirconium diboride by tungsten carbide additions, J. Am. Ceram. Soc. 91 (2008) 3530-3535. [18] V.A. Lavrenko, A.D. Panasyuk, T.C. Protsenko, V.P. Dyatel, E.S. Lugovskaya, E.I. Egorova, High temperature reactions of materials of the ZrB2-ZrSi2 system with oxygen, Porosh. Metall. 6 (1982) 56-58. [19] D. Sciti, M. Brach, A. Bellosi, Oxidation behavior of a presureless sintered ZrB2- MoSi2 ceramic composite, J. Mater. Res. 20 (2005) 922-930. [20] D. Sciti, M. Brach, A. Bellosi, Long term Oxidation behaviour and mechanical strength degradation of a pressureless sintered ZrB2-MoSi2 ceramic, Scr. Mater. 53 (2005) 1297-1302. [21] D. Sciti, V. Medri, L. Silvestroni, Oxidation behaviour of HfB2-15 vol.% TaSi2 at low, intermediate and high temperatures, Scr. Mater. 63 (2010) 601-604. [22] M. Kazemzadeh Dehdashti, W.G. Fahrenholtz, G.E. Hilmas, Effects of temperature and the incorporation of W on the oxidation of ZrB2 ceramics, Corros. Sci. 80 (2014) 221-228. [23] V.O. Lavrenko, A.D. Panasyuk, O.M. Grigorev, O.V. Koroteev, V.A. Kotenko, High-temperature oxidation of ZrB2-SiC and ZrB2-SiC-ZrSi2 ceramics 1700 °C in air, Powder Metall. Metal Ceram. 51 (2012) 217-221. [24] O.N. Grigoriev, B.A. Galanov, V.A. Lavrenko, A.D. Panasyuk, S.M. Ivanov, A.V. Koroteev, K.G. Nickel, Oxidation of ZrB2-SiC-ZrSi2 ceramics in oxygen, J. Eur. Ceram. Soc. 30 (2010) 2397-2405. [25] S. Guo, T. Mizuguchi, M. Ikegami, Y. Kagawa, Oxidation behavior of ZrB2- MoSi2-SiC composites in air at 1500 C, Ceram. Int. 37 (2011) 585-591. [26] H. Pastor, R. Meyer, An Investigation of the effect of additions of metal silicides on titanium and zirconium borides from the point of view of their sintering behavior and their resistance to oxidation at high temperature, Rev. Int. Htes Temp. Refract. 2 (1974) 41-54. [27] F. Peng, R.F. Speyer, Oxidation resistance of fully dense ZrB2 with SiC, TaB2, and TaSi2 additives, J. Am. Ceram. Soc. 91 (2008) 1489-1494. [28] E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2 and HfB2-based ultra-high temperature ceramics: effect of Ta additions, J. Mater. Sci. 39 (2004) 5969- 5977. I.G. Talmy, J.A. Zaykoski, M.M. Opeka, A.H. Smith, Properties of ceramics in the system ZrB2-Ta5Si3, J. Mater. Res. 21 (2006) 2593-2599. [30] S.C. Zhang, W.G. Fahrenholtz, G.E. Hilmas, Oxidation of ZrB2 and ZrB2-SiC ceramic with Tungsten additions, Electrochem. Soc. Trans. 16 (2009) 137-145. [31] D.G. Barton, S.L. Soled, G.D. Meitzner, G.A. Fuentes, E. Iglesia, Structural and catalytic characterization of solid acids based on zirconia modified by tungsten oxide, J. Catal. 181 (1999) 57-72. [32] L. Silvestroni, H.J. Kleebe, S. Lauterbach, M. Müller, D. Sciti, Transmission electron microscopy on Zrand Hf-borides with MoSi2 addition: densification mechanisms, J. Mater. Res. 25 (2010) 828-834. [33] H. Okamoto, The Zr-Si system, Bull. Alloy Phase Diagr. 11 (1990) 513-519. [34] L.Y. Chang, M.G. Scroger, B. Phillips, Condensed phase relations in the systems ZrO2-WO2-WO3 and HfO2-WO2-WO3, J. Am. Ceram. Soc. 50 (1967) 211-215. [35] H.A. Wriedt, The O-W (oxygen-tungsten) system, Bull. Alloy Phase Diagr. 10 (1989) 368-384. [36] A.B. Gokhale, G.J. Abbaschian, The Mo-Si system, 493-498. [37] L. Silvestroni, C. Melandri, S. Guicciardi, D. Sciti, TaB2-based ceramics: microstructure, mechanical properties and oxidation resistance, J. Eur. Ceram. Soc. 32 (2012) 97-105. [38] C.M. Carney, T.A. Parthasarathy, M.K. Cinibulk, Oxidation resistance of hafnium diboride ceramics with additions of silicon carbide and tungsten boride or tungsten carbide, J. Am. Ceram. Soc. 94 (2011) 2600-2607. [39] D. Caurant, O. Majérus, E. Fadel, M. Lenoir, C. Gervais, O. Pinet, Effect of molybdenum on the structure and on the crystallization of SiO2-Na2O-CaO- B2O3 glasses, J. Am. Ceram. Soc. 90 (2007) 774-783.  J. Phase Equilib. 12 (1991)  [29]  \\x0c']"
},{
  "_id": 149,
  "PDF": "Oxidation behavior of ZrB2-SiC laminates- Effect of composition on microstructure and mechanical strength.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 35 (2015) 1699-1714  Oxidation behavior of ZrB2 /SiC laminates: Effect of composition on microstructure and mechanical strength  E. Padovano a,∗  , C. Badini a , E. Celasco b , S. Biamino a , M. Pavese a , P. Fino a  a Politecnico di Torino, Department of Applied Science and Technology, C.so Duca degli Abruzzi 24, 10129 Torino, Italy b Università di Genova, Department of Physics, Via Dodecaneso 33, 16146 Genova, Italy  Received 1 August 2014; received in revised form 10 December 2014; accepted 26 December 2014  Available online 15 January 2015  Abstract  SiC-ZrB2 ceramic  laminates with different SiC:ZrB2 ratios were prepared by  tape casting, binder burnout and pressureless sintering at 2200 C in view of  their application  in  thermal protection systems. Their oxidation behavior  in air was  investigated by TGA up  to 1600 C and by  long term oxidation  tests at 1500 and 1600 C. Both microstructure and mechanical features of  laminates were compared, before and after oxidation. The formation of a passive layer was always observed but the oxidation resistance worsened and the complexity of the oxide layer microstructure increased with the ZrB2 content. Composition also affected the residual strength and modulus after oxidation. © 2015 Elsevier Ltd. All rights reserved.           Keywords: Oxidation resistance; Pressureless sintering; Tape casting; ZrB2 ; SiC  1.   Introduction     Both SiC  and  ultra-high  temperature  ceramics  (UHTCs) emerged as candidates for high temperature applications in the 1960s and, particularly during the past 15 years, they have been investigated for aerospace applications. Current TPSs comprising advanced SiC-coated Cf /C and Cf /SiC composites (CMCs) do not have  the oxidation resistance necessary at over 1650 C and cannot be reused for many missions. Nevertheless SiC can display passivating behavior up  to 1600 C, since a protective SiO2 layer  forms on  its surface  in  the oxidizing environment. The main  lack of  this coating deals with  its possible  failure caused by the thermal shocks and favored by the thermal expansion coefficient mismatch between SiO2 , SiC and CMCs.  In addition below 1600 C the operating conditions of space vehicles can also cause passive-to-active transition of SiC oxidation mechanisms. Actually the passive oxidation layer grown on SiC can be consumed or SiC can suffer recession when temperature greatly increases and oxygen partial pressure decreases contemporaneously. Anyway ground  tests performed under simulated        ∗  Corresponding author. Tel.: +39 0110904708.  E-mail address: elisa.padovano@polito.it (E. Padovano).  http://dx.doi.org/10.1016/j.jeurceramsoc.2014.12.029  0955-2219/© 2015 Elsevier Ltd. All rights reserved.        re-entry conditions  (oxygen partial pressure decreasing down to few mbar and  temperature  increasing up  to 1550 C during each  test)  indicated  that  the passive mechanism of SiC oxidation prevails over  the active one since ceramic  laminates made of pure SiC layers survived 100 re-entry cycles and the passive silica  layer still covered  the specimen surface at  the end of  the ground experiment.1,2 On  the other hand  the sharp profiles of space vehicles and hypersonic aircraft demand materials able to sustain temperatures even exceeding the decomposition temperature of SiC as well as  the melting  temperature of cristobalite (1725 C) which, after melting, evaporates. Among  the UHTCs considered for TPSs of hypersonic aircraft, ZrB2 has been  the most  investigated  since  it  exhibits a  unique  set  of  properties:  very  high melting  temperature (3245 C), good strength, modulus,  thermal shock  resistance, hardness and erosion  resistance. The best mechanical properties can be achieved only after  full sintering, and  they can be retained even at high  temperatures.3 On  the other hand ZrB2 has  some disadvantages:  the  rather poor  fracture  toughness characteristic of almost all ceramics, poor oxidation resistance, density of 6.085 g/cm3 (which  is quite high for TPSs) and  low intrinsic  sinterability. Because of  their poor  sinterability,  the ZrB2 -base ceramics have usually been processed by hot pressing at 1820-2000 C4-20 ; less frequently, these ceramics have been              \\x0c', '1700   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714     sintered by spark plasma synthesis21 or by pressureless sintering at 2150 C.22 In all cases the densification can be enhanced by using proper sintering aids; among  them SiC has proved  to be very effective. The addition of SiC can also improve flexure strength,4,5,8 KIc 5,17 and creep behavior.23 Oxidation  behavior  of ZrB2 and  some ZrB2 /SiC  composites  (containing  from  10  to  30 vol.%  of SiC)  has  been well  investigated  in  the  temperature  range  from  800  to 2000 recently well  reviewed by Fahrenholtz and Hilmas.20 Pure ZrB2 undergoes oxidation  starting from 800 C according to the reaction: ZrB2(s) +  C5,6,9-16,18,19,21 and    5/2O2(g) ↔   ZrO2(s) +   B2O3(s,l)  (1)              The  liquid  layer of B2O3 grants protection  from oxidation up to 1100-1200 C. Above 1200 C the volatilization of B2O3 , which exposes  the underlying porous ZrO2 scale, dramatically worsens  the oxidation resistance of ZrB2 .16,17,20 The addition of SiC particles  to a ZrB2 matrix greatly  improves oxidation behavior at high  temperatures since  it changes  the oxide scale microstructure.5,14,16-18,20,21 Below 1200 C ZrB2 /SiC composites behave  like pure ZrB2 since  they undergo oxidation more quickly than SiC thus giving a passive B2O3 glassy surface layer. Above this temperature also SiC quickly oxidizes according to the reaction: SiC(s) +   3/2 O2(g) ↔   SiO2(s) +   CO(g)  (2)              C).18,24  The resulting oxide scale obtained up to 1600 C was found to be very  similar under  every oxidizing  condition.5,9-21 It consists of an external borosilicate glassy layer and an underlying  layer made of a skeleton of crystalline ZrO2 , with voids filled  in by  the glassy phase. The B2O3 content within  the glassy phase decreases with  the  temperature  increase owing  to evaporation and at around 1500 C  the glassy  layer facing  the atmosphere mainly consists of silica. This external glassy layer should grant  improved oxidation protection since  the oxygen diffusivity  through pure silica  is  lower  than  that  through both −21 m2 /s for SiO2 , 10 −14 m2 /s for borosilicate and zirconia (10 −10 m2 /s for ZrO2 at 1500 B2O3 -21 mol.% SiO2 and 10 However at very high  temperatures also  the volatilization of molten silica occurs and the recession of the passive protective to silica.25 layer can prevail on  the oxidation of  the substrate  The external glassy  layer shows a complex structure resulting from the emission of gas bubbles and the outward flow of lowviscosity borosilicate liquid.24 In between the layer of ZrO2 plus borosilicate glass and the ZrB2 /SiC composite bulk some authors observed a porous zone of SiC depletion5,9,14,16,18,19 while others did not report any evidence of  this sub-scale.12,13,20,21 The SiC depletion is usually attributed to the active oxidation of SiC to gaseous SiO and CO caused by the low activity of oxygen in this zone.24 Since the diffusion of oxygen through the external layer of silica-rich glass occurs slowly and since the formation of silica results from the oxidation of SiC, according to some experimental results and modeling,26 oxidation resistance  improves with  the  increase of SiC content  in  the composite, at  least at temperatures up  to 1500-1600 C. On  the contrary, according to some  literature,27 when ZrB2-SiC composites are oxidized              at 1900 C it seems not advantageous to increase the percentage of SiC over 20%. Actually  the  formation of a porous  layer at the SiC depletion zone becomes more  important with  the  temperature  increase and  it could promote spallation of  the oxide layer.  In  the past only composites with a SiC content  ranging between 10 and 30 vol.% were  investigated, while very  little can be found  in  the most recent  literature about  the oxidation behavior of ZrB2 /SiC composites with high SiC percentage.18 From the above it can be inferred that at present there is not an ideal material to be used for oxidation protection in re-usable TPSs since such a kind of material  is expected  to operate  in a wide range of  temperature;  in fact SiC proved  to be suitable for this purpose at temperatures below 1600 C while ZrB2 -SiC composites are believed  to behave better at very high  temperatures (for  instance around 2200 C [28]) as well as under  low partial pressure of oxygen.29 On the other hand a laminate with alternating  layers of SiC and ZrB2 /SiC composite could display high oxidation resistance under a variety of environmental conditions. In addition ceramic  laminates showing weak  interfacial bonds between the layers can display improved toughness over monolithic ceramics because  the  interfaces and  the residual stresses cause crack deflection phenomena that increase the fracture work.30-34 Tape casting  is a powerful method  for  the production of green ceramic  tapes of whatever composition; after stacking  these  thin  tapes  the  laminate can be obtained by binder burnout and pressureless sintering by using proper sintering aids. The capability of producing multilayer ceramics (e.g. containing ZrB2 ) according to this method, and therefore without hot pressing,  should enable near-net  shape processing of complex components, thus reducing the cost with respect to hot pressing and machining conventional production methods. In view of developing hybrid multilayer ceramics  for TPS applications  the oxidation behavior of ceramic  laminates with different ZrB2 :SiC ratio should be deeply  investigated. In fact the  specific microstructure of ceramic  laminates  in  terms of interfaces between  the  layers could  result  in a different oxidation behavior with  respect  to  similar monolithic  ceramics produced by hot pressing. In this paper multilayer ceramic composites with composition (vol.%) ranging between 100%SiC and 20%SiC-80%ZrB2 as well as laminates consisting of alternate layers with different composition were prepared by  tape casting, binder burnout and sintering. Their oxidation behavior was investigated and both microstructure and mechanical behavior were compared before and after oxidation.  2. Materials and methods  2.1. Laminate preparation  The  laminates were produced  according  to  the  following processing path: preparation of a  slurry,  tape casting of  this slurry  for obtaining green ceramic  tapes,  tape drying by slow liquid evaporation, cutting  the  tapes and stacking  the ceramic layers according to proper sequences, binder burnout, final sintering performed without the aid of pressure. The slurries were produced by mixing  the ceramic powders,  liquids and additives in suitable proportions. Commercially available zirconium diboride powder  (Grade B, H.C. Starck, Germany, with  an  \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1701     average particle size of 2.25  \\u242em),  ␣-SiC powder (H.C. Starck UF-15, Germany, with a mean particle size of 0.55  \\u242em), C (Alfa Aesar flake 7-10  \\u242em) and B (H.C. Starck amorphous grade I, \\u242em) were used. The effect of C and B on the sintering mech1-2  anism and kinetics of SiC has been widely investigated.35-37 It is generally believed that during sintering C reacts with impurities containing oxygen (for instance with the SiO2 covering the surface of SiC particles) thus limiting the presence within the sintered ceramic of oxides that can melt at high temperatures.35,36 Boron is believed to act as an activator for SiC sintering that can also promote  the  formation of very  little amounts of B  Si C is over 0.5%.37 These sintering liquid when  its concentration  aids were adopted  in  this work because  the  formation at high temperatures of  liquid phases should be controlled and  limited for the specific TPS application, since these protection systems experience operating temperatures even exceeding 2000 C and have to keep suitable mechanical features also under this extreme condition. The ceramic powders made of SiC or mixtures of SiC and ZrB2 and sintering additives (C and B) were firstly dispersed for 12 h  in a mixture of ethanol, butanol and a dispersant (fish oil) by using an alumina jar with alumina milling bodies. Afterwards a plasticizer (polyethylenglycol, Bisoflex 102 Cognis) and a binder (polyvinilbutyral, Butvar B76 Solutia) were added  to the slurry  that was further mixed overnight. The ratio between solid and liquid components as well as the amount of binder and plasticizer were adjusted in order to obtain slurries with similar viscosities  irrespective of  the ratio between SiC and ZrB2 . To this purpose  the weight percentage of solid components of  the slurry ranged between 34.9% (SiC slurry) and 48.9% (20 vol.% SiC-80 vol.% ZrB2 slurry). For the same reason the percent of the binder was progressively decreased with the liquid percentage decrease from 9.6 to 7.6% while the amount of the plasticizer was decreased at the same time from 5.0 to 3.9%. The amount of sintering aids specifically added for promoting the SiC densification was calculated with respect to the SiC content in the slurry (1 and 3 wt.% of B and C respectively with respect to SiC). In the case of ZrB2-SiC composites the SiC itself acted as a sintering additive  that promotes  the ZrB2 densification. The  remaining part of  the slurries was a mixture of ethanol and butanol with a volume ratio of 4:6. The slurries were poured by tape casting on a mylar support moving with speed of 100 mm/min; doctor blade height of 1 mm. After  the  liquid evaporation at ambient temperature the green tapes (about 200  \\u242em thick) were detached form  the support, cut and stacked  in proper sequences of  ten layers or eleven  layers  (adopted  in  the case of  laminates with layers showing different composition with the aim of obtaining a symmetrical disposition of them). With the aim of improving the adhesion between  the green  layers, before stacking  them, their surface was wet with a glue made of ethanol, polyvinylalcohol and water. These  specimens were  then  submitted  to binder burnout under flowing argon atmosphere; this treatment, which consists  in a slow heating up  to 800 C, resulted  in  the thermal decomposition of organics (binder and plasticizer). The final  sintering  treatment was  carried out under  argon  atmosphere (99.99% purity) and controlled pressure below 550 mbar at 2200 C for 30 min (oven T.A.V. Cristalox, Italy). This  temperature  is needed  for obtaining good mechanical properties        ×  ×  and  relative density over 95% when pressureless sintering of SiC  is carried out by using C and B as sintering additives.38 Nevertheless also  the pressureless  sintering of ZrB2 must be conducted at  similar  temperature.22 In  this manner  laminate specimens with size 55 mm   12 mm   1.8 mm were prepared. Several kinds of laminates, differing one from another for the SiC and ZrB2 content but with all the layers showing the same composition, were produced. The percentages of SiC within each layer were: 100, 80, 60, 40 and 20%. Laminates showing alternate SiC and composite  layers were also  investigated. These laminates were designed according to the following architecture: 3SiC-1SZ-3SiC-1SZ-3SiC (where SiC and SZ stand for silicon carbide and composite layers respectively, while the figures stand for the number of layers of each sort stacked). Two kinds of these  laminates were prepared using either 20%SiC-80%ZrB2 or 50%SiC-50%ZrB2 composite layers.  2.2. Oxidation tests        ×  ×  ×  ×  Oxidation  resistance was  investigated by  submitting very small 3 mm   3 mm   1.8 mm specimens  (cut  from  the  larger sintered bars) to temperature runs under flowing air (50 ml/min) in a  thermal gravimetric analysis equipment  (Mettler-Toledo FTA-TGA). The samples were hold  in alumina crucibles. Two subsequent temperature runs comprising heating up to 1600 C (heating rate of 10 C/min) and cooling down to ambient temperature (cooling rate 18 C/min) were performed on each sample. The  behavior  of  laminates  containing  100,  80,  60,  40  and 20%SiC was compared. Also ceramic  laminates with alternate layers of SiC and SiC-ZrB2 composites were submitted to TGA experiment. The mass gain against  temperature was detected with  the aim of  investigating  the oxidation kinetics during  the first oxidation cycle and  the possible oxidation delay occurring during the second oxidation cycle because of the previous formation of a passive  layer. Some of  these  laminates  (bars 55 mm   12 mm   1.8 mm in size) were also submitted to longperiod oxidation treatments. This isothermal oxidation test was performed under  static air at  the  temperature of 1600 C  for 24 h in a Carbolite HTF1800 oven. A refractory frame with SiC fibers stretched between refractory supports was put  inside  the large furnace chamber. The specimens were hung by positioning  them on  the SiC fibers  in order  to avoid any contact with the walls of the oven. After long-time oxidation the morphology and  the microstructure of  the specimens were  investigated and compared  to  that of  the as-processed materials. Most of  these specimens  (markedly  those with  the higher content of ZrB2 ) swelled and  lost  their  regular  shape because of  the emission of gaseous oxidation products and  the  formation of borosilicate glassy phase  (with  low viscosity at 1600 C) during  this oxidation  treatment. For  this reason  it was not possible  to perform mechanical  tests on  these oxidized specimens according to current standards, which would be required for investigating the mechanical feature changes resulting from oxidation. In fact machining would be necessary in order to recover the geometrical shape needed for mechanical test, but machining would also remove  the external part of  the oxidized specimens and  therefore would change the overall composition of these specimens.           \\x0c', '1702   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714     Therefore less severe oxidation conditions were adopted in order to keep a suitable geometrical shape  for oxidized specimens. The 24 h  isothermal oxidation  treatment was  thus  repeated at 1500 C on the laminates less sensitive to oxidation owing to a limited content of ZrB2 (laminates with composition: 100%SiC, 80%SiC-20%ZrB2 and 60%SiC-40%ZrB2 ). The mechanical features of these last specimens were investigated.  2.3. Materials characterization  \\u242em   The microstructure and  the mechanical features of ceramic laminates under  investigation were compared before and after oxidation by  coupling  several  techniques. Their microstructure was  studied by X-ray diffraction  and micro-diffraction. Philips PW1710 equipment and Rigaku D/Max Rapid with a probe 100  in  size were used  respectively  for XRD  and \\u242e-XRD  analyses  (Cu K␣ radiation) of  laminate  surface  and selected zones of  laminate cross-sections. Optical microscopy and scanning electron microscopy coupled with energy dispersive spectroscopy (SEM-FEG Assing SUPRA 25; EDS Oxford) were used for checking microstructure, elemental analysis and elemental distribution maps on the laminate cross-section. The scale growing on ceramic  laminates was also studied by X-ray photoelectron spectroscopy (XPS PHI 5000 VersaProbe) with a monochromatic Al K␣  radiation. XPS analyses were performed by using an X-ray spot about 100  in diameter. Dedicated neutralization system was employed in order to avoid charging effect on  the samples during analyses.  It consists  in an electron gun, combined with positive Argon  ion gun,  in order  to increase his efficiency. Calibration of XPS equipment was performed by matching the literature binding energy values of Au  \\u242em   4f7/2, Cu 2p3/2 and Ag 3d5/2 peaks. Geometric density was calculated from specimen weight and dimensions (measured by a caliper) and compared to the theoretical one. Apparent density was also measured by using an automatic equipment (Quantachrome Instruments, Ultrapyc 1200e) exploiting the measure of the true sample volume by employing Archimede’s principle of fluid (He gas) displacement. Theoretical density was calculated from the content of SiC, ZrB2 , B and C within each slurry and taking  into account of  the carbonaceous material arising from decomposition of organics during the binder burnout treatment. This  last data was obtained by submitting  the organic components  to TGA  test under argon atmosphere. Elastic modulus was measured (according to ASTM E 1876) by using a method based on impulse excitation technique and analysis of transient natural vibration on Resonant Frequency and Damping Analysis MF basic  instrument  (IMCE n.v. Belgium). Three-point bending  test was performed adopting 40 mm outer span and a cross-head speed of 0.1 mm/min (Sintech 10 D equipment, UNI EN 658.3 standard). Vickers microhardness measurements were performed on the cross section of the specimens with a test load of 300 g and a dwell time of about 10 s.  3. Results  3.1. Laminates with all layers showing the same composition  3.1.1. As-processed laminates  In every case  the microstructure of  the as-fabricated  laminates  consists of  several phases  (Fig. 1  and Table 1). The  Fig. 1. Microstructure of ceramic laminates (consisting of layers with the same composition) observed by SEM (SE): (a) SiC laminate, A* = SiC microstructure after  C, (b) 60%SiC-40%ZrB2 laminate; (c) 20%SiC--80%ZrB2 laminate.     thermal etching at 1600  Table 1  Analysis of particle grain size in function of the laminates chemical composition.  Material  composition  Phases   Apparent density (%  of theoretical)  100%SiC   SiC, C, B   80%SiC-20%ZrB2  SiC, ZrB2 , C, B   ±   0.2   94.8   ±   0.2   95.6   60%SiC-40%ZrB2 40%SiC-60%ZrB2 20%SiC-80%ZrB2  SiC, ZrB2 , C, B  ZrB2 , SiC, C, B  ZrB2 , SiC, C, (B)   99.0   93.8   99.3   ± ± ±   0.2    0.1    0.1   Grain size  SiC   Polygonal (few) 1-4   \\u242em  Rod like (many): length 5-20   \\u242em  aspect ratio: 3-7  Polygonal (many): 2-20   Rod like (few): length 5-20   \\u242em \\u242em  aspect ratio: 3-7  Polygonal: 2-20   Polygonal: 2-20   Polygonal: 2-20   \\u242em  \\u242em  \\u242em   ZrB2  -   C  5-10   \\u242em  Polygonal: 2-10   \\u242em   5-10   \\u242em  Polygonal: 2-10   Polygonal: 2-10   Polygonal: 2-10   \\u242em  \\u242em  \\u242em   5-10   5-10   5-10   \\u242em \\u242em \\u242em    \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1703  microstructure of SiC laminate comprises SiC grains and carbon inclusions while that of SiC-ZrB2 composite laminates consists of a homogeneous mixture of  the crystal grains of  these  two components, in addition also carbon inclusions can be observed (Fig. 1). Carbon flakes have been added as sintering aid and further carbon comes from the thermal decomposition of binder and plasticizer occurring during  the binder burnout  treatment. The tight mixture between phases rich in Si and Zr respectively is still kept after oxidation in the inner part of the specimens while the Si to Zr atomic ratio in the external oxide layer (growing outwards) changes. This is due to convective motions of a glassy phase rich in silicon and it depends also on the diffusion of these elements (Si, Zr) toward the surface and their different tendency to react with oxygen, as will be better discussed in the following sections. Most of  the carbide crystals  in  the SiC  laminate show rod-like shape and are mixed with smaller polygonal grains (Fig. 1A*). The length of these elongated crystals generally ranged between five and  twenty  \\u242em (but crystals as  long as 100  \\u242em were also observed) and  the size of  the polygonal ones was around  few micrometers, therefore noticeable grain growth of SiC occurred during sintering. The main components of SiC-ZrB2 composites were polygonal mixed grains of these two ceramics with size ranging from 2 to 20  \\u242em. Therefore in the case of the composite laminates  lower grain growth occurred during sintering. X-ray diffraction confirmed  that  the  laminates contained SiC, ZrB2 and C (Fig. 2a-c). It was not possible to distinguish clearly the boundaries between adjacent  layers on  the cross section of  the single composition laminates, however they can be appreciated on  the  fracture surface where  the  layers are separated owing to delamination phenomena, as previously reported.39 This feature  is caused by a certain weakness of  the  interfacial bonds between neighboring  layers, which  is a characteristic of  these laminates.  3.1.2. TGA oxidation tests        The TGA results are summarized in Fig. 3. The TGA curve of SiC  laminate  (Fig. 3a) shows only very small weight gain starting from 800 C (less than 0.3%) according to the oxidation reaction of SiC  (reaction 2); starting  from about 1500 C  the mass gain  is more  than counterbalanced by  the oxidation of carbon  to gaseous carbon oxide. This  thermogram shows  that only  little SiC undergoes oxidation and  therefore  that a  thin silica layer grows on the sample surface. On the other hand this thin silica layer grants good protection against further oxidation since the repetition of the TGA run under flowing air on the same sample results in a flat curve (total mass gain at 1600 C lower than 0.2%). The laminate containing 80 vol.% of ZrB2 (Fig. 3d) displays worse oxidation resistance, in fact the first run in air up to 1600 C results  in about 7% mass gain. The weight starts  to increase at about 800 C owing the oxidation reactions 1 and 2, but  the oxidation rate further  increases starting from 1400 C. Nevertheless the oxide layer grown on the sample surface during the first TGA run still displays passive effect since during  the second run under flowing air the mass of the 80%ZrB2 -20%SiC laminate increases of less that 2% only (Fig. 3d). The composite  laminates containing 20 vol.% of ZrB2 and 40 vol.% of ZrB2 show mass gain of 3 and 5.7%  respectively              Fig. 2. X-ray diffraction  (XRD) spectra of  laminate surface and micro-X-ray diffraction (\\u242e-XRD) of laminate cross-section: (a) XRD of SiC laminate surface;  (b) XRD of 60%SiC-40%ZrB2 laminate surface; (c) XRD of 40%SiC-60%ZrB2 laminate  surface;  (d) XRD of SiC  laminate  surface after oxidation  in air at \\u242e-XRD of SiC laminate cross section after oxidation in air at \\u242e-XRD of 60%SiC-40%ZrB2 section after oxidation, glassy external  layer;  (g)  \\u242e-XRD of  C for 24 h, zone below the passive layer; (f)   laminate cross   C for 24 h; (e)   1600  1600       60%SiC-40%ZrB2 laminate cross section after oxidation,  intermediate oxide layer; (h)  \\u242e-XRD of 60%SiC-40%ZrB2 laminate cross section after oxidation, inner porous layer.  during the first TGA cycle under air and the scales that form on the surface are able  to reduce  the oxidation occurring during a second oxidation cycle (Fig. 3b and 3c). Conclusively the oxidation of all the laminates results in the formation of a passive layer, whose effectiveness increases with the increase of the SiC content in the laminate.  \\x0c', '1704   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714  Fig. 3. TGA oxidation test under flowing air up to 1600     C (repeated twice on the same sample; continuous line = first oxidation cycle, dashed line = second oxidation  cycle): (a) 100%SiC laminate; (b) 80%SiC-20%ZrB2 laminate; (c) 60%SiC-40%ZrB2 laminate; (d) 20%SiC-80%ZrB2 laminate; (e) laminate with alternate layers of 100%SiC and of 20%SiC-80%ZrB2 composite.  3.1.3. Long-term oxidation tests     The ceramic  laminates under  investigation behave  in very different manners when  submitted  to  isothermal oxidation at 1600 C under static air. The characteristics of  the oxide  layers grown during  the 24 h  treatment under  these conditions are compared in Fig. 4 and Table 2. The thickness of the oxide layer increases with  the  increase of  the ZrB2 content. The  laminates containing 100 and 80 vol.% of SiC show a continuous but rather thin oxide  layer (Fig. 4a and b), while  the specimens containing at  least 40 vol.% of ZrB2 display  the formation of surface oxide layers with complex microstructure which consist of three parts: an external glassy  layer, an  intermediate oxide  layer and a porous  layer  (Fig. 4c, d and e). The  specimens with more than 40 vol.%ZrB2 blistered during oxidation and  this  feature resulted in an external layer with irregular thickness and showing the presence of bubbles inside, very likely caused by the emission of gaseous phases. Actually the oxidation of SiC entails the formation of carbon oxides and also  the boron oxide resulting from  the oxidation of ZrB2 melts and volatilizes at  temperatures over 800 C. The  long  term oxidation of specimens with the highest ZrB2 content (namely 80% and 60 vol.%) caused an     appreciable  increase of  the overall sample  thickness, which  is consistent with the formation of a scale growing outwards. Also the thickness of the intermediate oxide layer increases when the ZrB2 percentage is very high (80 vol.%) while the thickness of the porous layer seems scarcely affected by the ZrB2-SiC ratio in the  laminate (Table 2 and Fig. 4). However from  the morphology  it could be  inferred  that some differences  in  the porosity degree exist for specimens with different composition (Fig. 5c and  f). Also  the morphology of  the  intermediate oxide  layer changed with the laminate composition (Fig. 5b and e). This part of the scale always consisted of light gray crystals embedded in a dark matrix; these two phases give rise to a columnar structure grown perpendicularly to the surface in specimens with 80 vol.% of ZrB2 (Fig. 5b) while  interconnected networks of both  these phases were observed in the laminates with a lower ZrB2 content (Fig. 5e). Accurate analyses performed by using XRD, micro-XRD, EDS  and XPS  techniques  allowed  to  better  investigate  the microstructure of  the oxide scale. The XRD of  the surface of the  laminates containing 80% and 100 vol.% of SiC shows  the presence of cristobalite, SiC and graphite after oxidation. The  Table 2  Thickness of oxide layers grown on the surface of laminates with different composition after oxidation treatment under static air for 24 h at 1600  Material composition   Layer thickness (\\u242em)  External glassy layer   Oxide intermediate layer   Porous layer   100%SiC   80%SiC-20%ZrB2 60%SiC-40%ZrB2 40%SiC-60%ZrB2 20%SiC-80%ZrB2  -   -   30-164   83-128   600-700   aOxide scale made of a single layer mainly containing silica.  21-36a 97-115a  192-256   105-180   1200-1300   -   -   95-170   120-230   140-200      C.  Sample core  1750  1700  750-1200  1300-1600  640  \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1705  Fig. 4. Oxide   layer grown on   the   laminates with different composition during 24 h oxidation   in air at 1600     C   (sample cross-sections):   (a) SiC   laminate;   (b)  80%SiC-20%ZrB2 (c) 60%SiC-40%ZrB2 (d) 40%SiC-60%ZrB2 intermediate oxide layer and sample core in 20%SiC-80%ZrB2 laminate.  laminate;   laminate;   laminate;   (e) 20%SiC-80%ZrB2  laminate;   (f) porous   layer   in between  Table 3  EDS analyses on the different parts constituting the oxide layer grown on laminates with different composition after oxidation in static air for 24 h at 1600  Sample   Zone under investigation   Atomicpercentage of elements  100%SiC  Passive layer   Below the passive layer   External glassy layer   60%SiC-40%ZrB2  Dark area in the intermediate zone   Light gray crystals in the intermediate zone   Porous layer   External glassy layer   20%SiC-80%ZrB2  Dark area in the intermediate zone   Light gray crystals in the intermediate zone   Porous layer   Si   29.89   42.66   25.86   22.56   -   <0.06   14.76   20.11   -   1.51   O   66.73   2.76   61.66   64.42   56.94   6.05   61.46   69.71   60.13   28.74   C   3.38   54.58   1.38   0.38   9.58   15.63   17.13   2.08   8.69   40.47   Zr   -   -   0.17   0.07   30.51   29.37   0.56   0.21   26.97   8.55      C.  B  -  -  10.92  12.58  2.93  48.84  18.44  7.88  4.21  20.82  \\x0c', '1706   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714  Fig. 5. Cross sections of   C. 20%SiC-80%ZrB2 laminate: (a) overview of  with rod-like ZrO2 , (c) porous layer. 60%SiC-40%ZrB2 laminate: (d) overview of the oxide layer, (e) intermediate layer with interpenetrated ZrO2 and glassy phase networks, (f) porous layer.  laminates after oxidation   in air for 24 h at 1600     the oxide   layer, (b)   intermediate   layer  XRD spectrum of oxidized SiC laminate is shown in Fig. 2d, a similar spectrum was observed after oxidation for the laminate with composition 80%SiC-20%ZrB2 . Silica  is  the main component while the intensities of peaks belonging to both SiC and C greatly decreases with respect  to  the XRD pattern of  the asfabricated  laminate (Fig. 2a). The EDS analysis carried out on the sample cross-section confirms  that  the oxide scale mainly consists of  silica  (O:Si atomic  ratio close  to 2:1) with  some inclusions of  residual carbon  (Table 3). Just below  the oxide layer some oxygen diffused through the passive layer was found in the matrix which mainly consists of SiC (Table 3 and Fig. 2e). The three sub-layers constituting the scale grown on all the laminates containing from 40%  to 80 vol.% of ZrB2 showed very different compositions (Table 3). Table 3 shows the composition of sub-layers obtained for specimens with the limit compositions of this range. The main elements contained within the external glassy layer are O, Si and B; some carbon was also found while the content of Zr was very  low. The atomic  ratio between O, Si and B  is consistent with  the formation of silica and boria. It must be underlined  that not negligible amount of boron  is still present here after 24 h of oxidation in air at 1600 C and the percentage of boron found here  increases with  the  increase of  the ZrB2 percentage at the sample core. This means that boria evaporation is counterbalanced by boron or boria diffusing from the inner part of the specimen, on the other hand at a visual inspection  the sample surface after  the oxidation  treatment appears covered by a glassy  layer. X-ray micro-diffraction of  this zone shows only  few and weak peaks belonging  to ZrO2 superimposed  to broad signals  that are characteristics of glassy phases (Fig. 2f). XPS results allow  to confirm  the formation of silica since  there  is a clear evidence of Si O bonds  in  the Si2p and O1s spectra  (Fig. 6a). The C1s spectrum shows mainly C C bonds while only little carbon is bonded to Si or O, therefore in this area  inclusions of graphite, partially oxidized, are present (Fig. 6a). Unfortunately the low sensitivity of XPS with respect to B does not allow to obtain information about the bonds formed     by boron. As a matter of fact the XPS sensitivity factor of boron (0.15 only) is much lower than those of carbon (0.29) or oxygen (0.71). All these results bring to the conclusion that the external layer consists of borosilicate glass with graphite inclusions and large pores. The intermediate oxide layer comprises two different phases whose chemical composition detected by EDS  is summarized in Table 3 for 40%ZrB2-60%SiC and 80%ZrB2-20%SiC composites, taken as representative examples. The phase appearing darker at SEM observation mainly contains O, Si and B, while the atomic percentages of both C and Zr are here very low. For these laminates the atomic ratio between oxygen and silicon plus boron detected within the dark zones is consistent with the presence of SiO2 plus B2O3 , that is with borosilicate glass. The light gray phase mainly contains Zr and O  instead, but not negligible amount of C and B (about 9 at.% C and 3-4 at.% B) were also observed. Using EDS, micro-XRD and XPS  techniques  it is hard  to obtain  information related  to each single component of this two-phase oxide layer, because of the limit of spatial resolution of  these  techniques (respect  to  the single extension of some phases) and  the consequent difficulty  to make sure  that the beam  focused on one phase only. Nevertheless  the  results obtained with different  techniques were consistent. The main crystalline component  in  the XRD spectrum of  this area was ZrO2 ;  in addition  low-intensity peaks belonging  to ZrB2 and SiC were found, showing that the oxidation of the former components of the laminate did not occur completely here (Fig. 2g). The XPS spectra of Si2p, Zr3d and O1s show that Zr and Si are mainly combined with oxygen (Fig. 6b). The  lack of  the silica peak in the relevant XRD pattern indicates that glassy borosilicate forms. Therefore analytical techniques concur to show that in every case the dark zones belong to borosilicate glassy phase and the light gray ones to ZrO2 . On the other hand the morphology of the intermediate layer changes with the ZrB2 percentage in  the  laminate submitted  to oxidation. For  the  laminates containing 80 vol.% of ZrB2 columnar glassy zones placed parallel  \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1707  Fig. 6. XPS analyses performed on the different parts of the scale grown after oxidation in air (24 h, 1600  C) on a 80%ZrB2 -20%SiC laminate: (a) external glassy oxide layer (Si2p, O1s, C1s spectra); (b) intermediate oxide layer (Si2p, O1s, Zr3d spectra); (c) porous layer (Si2p, Zr3d, C1s spectra); (d) cross section core (Si2p,     Zr3d, O1s spectra).  to  the direction of oxygen diffusion can be observed, as previously reported  in  the  literature.12,40,41 When  the ZrB2 content decreases down to 60 or 40 vol.% the morphology of the intermediate layer changes and the elemental maps across the interface  between  the external and  the  intermediate  layer show  that Si and O concentrates  in  the glassy phase, which  is a continuous layer on  the  surface but gives  rise  to a network without any preferred orientation and interpenetrated with a complementary  \\x0c', '1708   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714  Fig. 7. Cross section of 60%SiC-40%ZrB2 laminate after oxidation in air (24 h, 1600 intermediate layer, (c) elemental maps of Zr, C, Si and O.     C): (a) overview of the oxide layer, (b) interface between the external and the  the   intermediate   layer  two-phase   rich of Zr within   network  (Fig. 7). A porous layer was found at the interface with the sample core in all oxidized specimens containing at  least 40% of ZrB2 . In addition the porous layer thickness was found not to be related to  the ZrB2 :SiC  ratio  in  the ceramic  laminate and  then  to  the morphology of  the  intermediate  layer  for specimens containing from 40  to 80% of ZrB2 . The micro-XRD patterns of  this zone show that the main crystalline phase is ZrO2 but also very weak peaks belonging to ZrB2 and residual SiC can be observed (Fig. 2h). The EDS analysis shows  that  the silicon content  is very  low  (less  than 2 at.%) but  the content of carbon  ranges between 16 and 40 at.% depending on  the kind of ZrB2 -SiC laminate (Table 3). The elemental maps  taken across  the  interface between the porous layer and the sample core (Fig. 8) show  that the porous layer is a zone of silicon depletion but not of carbon depletion; in addition also zirconium and oxygen are present within the porous layer. XPS spectrum for C1s demonstrate that chiefly C C bonds exists here, while only little carbon is bonded to O and Si (Fig. 6c). All these data bring to the conclusion that SiC  is consumed by oxidation and  this results  in  the depletion of silicon and  formation of amorphous carbon. As shown by XRD and XPS  (Figs. 2h and 6c)  the silicon still  remaining  is mainly combined with oxygen and only little SiC can be found; in addition Zr O bonds were detected by XPS. The long-term oxidation test confirmed the TGA results since all  the  analytical  results brought  to  the  conclusion  that  the oxidation  resistance under air  in  the  temperature  range up  to 1600 C of ZrB2-SiC  laminates progressively gets worse with the increase of the percent of ZrB2 in the material.     \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1709  Fig. 8. Element distribution at   long-term oxidation: (a) 20%SiC-80%ZrB2 laminate (porous  right and upper part of the maps; elemental maps of Zr, B, C, Si, O); (b) 40%SiC-60%ZrB2 laminate (porous layer in the right part of the maps; elemental maps of Zr, B, C, Si, O).  layer   the  the   interface between porous   layer and sample core after   in   3.2. Laminates with alternate layers  The microstructure  characteristic of  laminates  integrating composite  layers  in between  layers of SiC  is shown  in Fig. 9. Two kinds of specimens were prepared with the following architecture: 3SiC-1SZ-3SiC-1SZ-3SiC (were SiC and SZ stay for  pure SiC  layer and composite  layer  respectively). The  laminates with composite layers containing either 50 vol.% of ZrB2 or 80 vol.% of ZrB2 were successfully prepared according  to the method used  for  the  laminates having all  layers with  the same composition. They showed very similar microstructure. Some cracks were observed inside the composite layers mainly  \\x0c', '1710   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714  3.3. Mechanical and physical properties of laminates  The properties of  the  laminates are compared  in Table 4. Density values (geometrical) ranging between 90 and 93% of the theoretical ones were observed in every case for the as-processed materials with layers showing a single composition. The rather low geometrical density mainly is due to the surface roughness of the as-sintered materials (which greatly affects the measurement of  the volume of  these rather  thin  laminates) but  it also could result from residual micro-porosity, likely present at the interface between adjacent  layers. Actually  the apparent density values measured by using  the pycnometer were  significantly higher (between 94 and 99% of theoretical density) than the geometrical ones, thus confirming that geometrical density values are mainly affected by the roughness of the surface of the specimens. Elastic modulus progressively increased and the bending strength only slightly decreased with the increase of the percentage of ZrB2 , which is consistent with the mechanical features of pure SiC and ZrB2 .29,42-44 A clear trend of micro-hardness with the change of composition could not be appreciated. The composite laminates with different composition also show some differences in terms of size and shape of crystal grains and porosity. In principle also these  features could affect  the mechanical behavior. The density of  laminates containing at  least 80% of SiC  is a bit  lower than that of laminates containing 60 or 20% of SiC only, but the higher residual porosity of laminates containing at least 80%SiC did not result  in  the decrease of bending strength, while  these materials showed  lower modulus  than other  laminates. Porosity should affect negatively both strength and stiffness, but this was not observed. Polygonal crystal grains of similar size were always observed in composite laminates, but also larger rod-like grains of SiC were detected  in  the microstructure of  laminates containing at  least 80% of SiC. The comparison of specimens with  the  same  relative density and  similar  size and  shape of crystal grains but different composition shows that mainly chemical composition is responsible for the mechanical behavior. For instance the laminate with composition 20%SiC-80%ZrB2 has the same density and microstructure of  the 60%SiC-40%ZrB2 but very different strength and stiffness  (Table 4). Also when comparing the mechanical characteristics of laminates with different size and shape of crystal grains  it cannot be  found any clear effect of this morphological features. For instance samples containing elongated and large grains (like 100%SiC laminate) show higher strength but  lower modulus with  respect  to  laminates with polygonal  crystals only  (like 20%SiC-80%ZrB2 laminate). Conclusively the mechanical behavior of these laminates was found to depend chiefly on the SiC:ZrB2 ratio, because of the different intrinsic mechanical characteristics of these two ceramics. In order  to  investigate  the effect of  the oxidation process on  the mechanical characteristics, ceramic  laminates containing less than 50 vol.% of ZrB2 were mechanically tested before and after oxidation  in static air at 1500 C  for 24 h. This oxidation  treatment  slightly decreased both elastic modulus and strength of SiC  laminate. The same oxidation  treatment more heavily worsened  the density and  the mechanical behavior of composite laminates containing ZrB2 . The geometrical density     Fig. 9. Microstructure of   laminate with 20%SiC-80%ZrB2 composite  (light gray) put between SiC  layers (dark gray); cross section showing cracks  layers  within the composite layer.  ×        ×  C).42,43  oriented perpendicularly  to  the  laminate  thickness  (Fig. 9). These defects can be attributed to the thermal mismatch between the composite and SiC as well as  to  the difference  in shrinkage between SiC and composites containing ZrB2 which occurs during binder burnout and sintering. Actually the thermal expansion  coefficient of ZrB2 is  from 7  to 27% higher  than  that of SiC depending on  the  temperature considered (e.g. CTE of −6 K −1 and CTE of SiC = 5.5  −6 K −1 , in the ZrB2 = 5.9   10  10 temperature range from ambient  temperature  to 1000 In addition  the overall  shrinkage occurring during  the whole material processing (organics burnout plus sintering treatments) was 42.3 vol.% for SiC and 46.8 vol.% for 80%ZrB2 -20%SiC multilayer composite, but major differences  in shrinkage were observed after organics burnout (4.0 and 9.8 vol.% for SiC and composite respectively). The different shrinkage of these materials during the thermal treatments can be attributed to both the different composition and  the different size of  the powders of SiC and ZrB2 . During  thermal  treatments  the composite  layers suffer major shrinkage and during cooling from the sintering temperature the composite layers are affected by greater thermal contraction. These phenomena cause tensile stresses in the plane of the composite tapes which are responsible for the crack formation. The TGA oxidation test performed in air up to 1600 C showed  that  the oxidation behavior of  the  laminate  integrating composite layers (Fig. 3e) is very similar to that of SiC laminate and much better  than  those of composite  laminate containing 80 vol.%. of ZrB2 within every layer (Fig. 3d). Actually most of  the  sample  surface  facing  the oxidizing atmosphere  is constituted by SiC while only a  little part of  it consists of SiC-ZrB2 composite, since the composite layers are exposed  to air on  the  lateral  side of  specimen only. For  this reason  the mass gain during  the first TGA  run was about 1% only and no mass gain was observed after the second oxidation test performed on the same specimen. Conclusively these TGA results proved that the external SiC layers that undergo passivation are able  to grant oxidation protection  to  the composite layers containing ZrB2 .  \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1711  Table 4  Physical properties of laminates before and after long-term oxidation treatment.  Sample   Geometrical density  (% of theoretical)  Apparent density  (% of theoretical)  Elastic modulus  (GPa)  Bending  strength (MPa)  Microhardness  (GPa)  100%SiC   89.0   ± ± ± ± ± ± ± ± ± ±   2.6  94.8   ± ± ± ± ± ± ± ± ± ±   0.2  339   ± ± ± ± ± ± ± ± ± ±   19  324   ± ± ± ± ± ± ± ± ± ±   24  21.2   ± ± ± ± ± ± ± ±   0.6  100%SiC after oxidation at 1500     C  89.7    1.0  95.9    0.2  336    15   274    17   20.7    2.1  80%SiC-20%ZrB2 80%SiC-20%ZrB2 after oxidation at 1500 60%SiC-40%ZrB2 60%SiC-40%ZrB2 after oxidation at 1500 40%SiC-60%ZrB2 20%SiC-80%ZrB2 Alternate layers: SiC and 20%SiC-80%ZrB2 Alternate layers: SiC and 50%SiC-50%ZrB2  90.2    2.9   95.6    0.2   364    19   361    34   25.5    0.8     C   81.1    2.5   94.1    0.2   127    11   196    5   22.4    2.4  92.3    0.9   99.0    0.2   389    11   325    41   24.4    1.0     C   61.1    3.7   88.6    0.1   164    29   176    28   20.8    2.2  91.9    1.3   93.8    0.1   408    16   284    19   23.0    0.7  92.9    1.0   99.3    0.1   444    10   277    29    1.5  69.6    5.7   87.6    1.2   209    23.0   109    12.1   -  71.9    5.1   86.6    2.5   235    27.0   179    12.1   -  greatly decreased (up to about 30%) particularly in the case of the laminate with the higher ZrB2 content. Much more limited density variations were observed by using the Archimede’s method, which is less affected by the features of the sample surface. Both modulus and strength of composite laminates were appreciably lowered by this prolonged oxidation, but the final stiffness and strength of the oxidized composite laminates still seem sufficient for application  in TPS, since for  this application high mechanical performance  is not requested while oxidation resistance  is the main material requirement. The sintered laminates integrating composite layers in between SiC sheets showed geometrical and apparent density after sintering corresponding respectively to about 75 and 87% of  the  theoretical one. The presence of defects  (micro-cracks) and  therefore  the not  full densification observed on  laminates containing  two composite  layers out of eleven also heavily worsened the mechanical behavior (Table 4). In addition the presence of cracks in these laminates did not allow to perform affordable micro-hardness measurements.  4. Discussion  The composition of  the  laminates greatly affected  the oxidation behavior  in  terms of:  thickness and microstructure of the oxide layer, its capability to provide passivation and residual mechanical properties retained after severe oxidation. The thickness of the oxide scale dramatically increased with the increase of the ZrB2 /SiC ratio in the laminate. Nevertheless also the thickness of  the oxide  layer grown after  long-term oxidation on  the surface of  laminates fully made of SiC or mainly made of SiC (namely  laminates with 80 vol.% of SiC)  seems  rather high. Actually  it was  found  to be much higher  than  that observed for high-purity monolithic SiC oxidized in similar conditions.45 Anyway the SiC laminate is well far to be pure SiC since it contains carbon, boron (sintering additives) and the solid residue left by organics (binder and plasticizers) after  the burnout. Therefore it seems that impurities greatly affect the oxidation kinetics. In addition the external layer facing the atmosphere after severe oxidation still contained inclusions of carbon embedded in crystalline silica or  in glassy borosilicate, notwithstanding carbon easily can undergoes oxidation. The TGA results confirmed the detrimental effect of ZrB2 on  the oxidation  resistance at  temperatures up  to 1600 C. According  to  the  reaction Eqs.  (1) and  (2),  the oxidation of SiC  to  silica entails 49.8% weight     increase while oxidation of ZrB2 to zirconia and boria results in 70.9% weight increase, but this last weight increase becomes 9.2% only  if boria evaporates. The TGA  results showed  that the weight gain of  the samples dramatically  increases with  the increase of  the ZrB2 percentage  in  the  specimen  in  spite of boria evaporation. This means that the fraction of laminate that undergoes oxidation increases with the ZrB2 content in the specimen, which  is well consistent with  the  thickness of  the oxide scales measured for specimens with compositions ranging from 100%SiC  to 20%SiC-80%ZrB2 . In addition  the passive effect of the cristobalite layer (grown on 100%SiC laminate) was better  than  that displayed by  the oxide  layers consisting of  three sub-scales (grown on ZrB2-SiC composite  laminates), as outlined by  the TGA experiments carried out on specimens  that previously suffered a first oxidation run. These outcomes also show  that  the effect of composition on oxidation resistance of laminates  is similar  to  that observed for monolithic SiC-ZrB2 ceramic composites.18 The shape of TGA curves obtained during the second oxidation run shows that the passive layer formed on  the surface of composite  laminates becomes  less effective starting  from 1400 C. As a matter of  fact  the volatilization of boria resulting from ZrB2 oxidation becomes progressively more  important with  the  temperature  increase;  for  instance according  to some  literature during  isothermal oxidation over 1200 C boria evaporation rate  increases so much  that boria  is no more present in the external part of the oxide layer after some time.9,20 During  the TGA run  the oxidation kinetics  increases too when  the volatilization of boria occurs very quickly. Very likely at high  temperature  the amount of glassy phase within the scale decreases with  respect  to  that of zirconium dioxide owing to the boria volatilization and, consequently, the passive effectiveness of  the  scale changes  (according  to  the  fact  that −14 m2 /s the oxygen diffusion coefficient at 1500 C is 1.7   10 −10 m2 /s for ZrO2 [18,24]). In for  the borosilicate glass and 10 addition the continuity of the passive layer is interrupted by the emission of gaseous species. In spite of the evaporation of boria not negligible amount of boron was found in the glassy layer even after  long-term oxidation probably owing  to boron diffusion toward the sample surface. In principle also the residual porosity should decrease  the oxidation  resistance of  the ceramic  laminates under investigation. Nevertheless, in this case, the effect of porosity is negligible since the sample surface is quickly covered by a passive layer during the first step of the oxidation process.           ×  \\x0c', '1712   E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714     In fact  the  laminate containing 100%SiC, which shows higher porosity than most of the composite laminates, displays the best oxidation resistance, very likely because of the superior passive effect of the silica passive layer. In addition SiC-ZrB2 laminates with very similar densities but different SiC:ZrB2 ratios showed very different oxidation resistance. Very different oxide  scale microstructures were observed depending on the laminate composition. In particular the whole layer mainly consisted of cristobalite when  the  laminate contained at  least 80 vol.% of SiC while an oxide scale consisting of  three sub-scales (external borosilicate glassy  layer,  intermediate two-phase layer, internal porous layer) was observed when the ZrB2 content in the laminate was equal or exceeded 40 vol.% It was not found any evidence for zircon formation, which can result from a reaction between zirconia and silica as previously reported by Gao et al.46 The formation of a complex oxide scale is consistent with most of the literature concerning high temperature oxidation in air (at 1500 C or more) of ZrB2 -based ceramics containing from 70  to 86 vol.% of ZrB2 ,7,9,20 even  though  the formation of the porous layer during oxidation at high temperature of ZrB2 -SiC composites with ZrB2 content of 80% or more was not observed by several authors.13,21,27 According to these contradictory outcomes  the  formation of a porous  layer close to  the not-oxidized sample core does not seem  related  to  the ZrB2-SiC ratio  in  the material submitted  to oxidation. On  the contrary Williams et al.18 proposed  that a correlation between the composite composition and the formation of the porous layer exists, since they found that the porous layer does not form when the ZrB2 content is lower than 50 vol.%. In the present paper the porous  layer was detected after oxidation of all  the  laminates containing at least 40 vol.% of ZrB2 and also the microstructure of the intermediate sub-scale only partially agreed with the literature. Therefore laminates behaved differently from monolithic ceramics. The morphology of  the  intermediate  two-phase  layer was characterized by a columnar structure grown perpendicularly to the surface in specimens with 80 vol.% of ZrB2 while interconnected networks of  the  two phases appeared  in  the  laminates with a  lower ZrB2 content. The columnar  structure, already described in some literature,12,40,41 probably inspired the modeling of the oxidation process for ZrB2 -based composites; in fact it has been proposed that the glassy phase allows for convection movements of glass and  therefore creates a  localized  inward path for oxygen.24 In addition, according to Parthasarathy et al. model,25 the oxygen  transported  in such a way  is  responsible for  the active oxidation of SiC, which results  in  the formation of the porous layer at the interface with the sample core. In this paper  the morphology of  the  intermediate  layer was not found to be completely consistent with this model when the ZrB2 content decreases down  to 60 or 40 vol.%. In  these specimens  the straight channels for oxygen diffusion are  lost and  the convective motions of liquid glass should be hindered by the network of ZrO2. This modification of the oxide layer morphology, and perhaps of  the diffusion mechanism as well,  is consistent with the decrease of  the overall  thickness of  the oxide  layer, which progressively occurs when  the ZrB2 content decreases  in  the ceramics.  The formation of  the porous  layer  is generally attributed  to SiC depletion  in  literature.7,9,18-20,25,26,46 The activity of  the oxygen diffusing through the scale, that can reach even the sample core  (see Fig. 6d),  reasonably decreases with  the distance from the sample surface and this feature causes active oxidation of SiC close  to  the  sample core. The conditions of  temperature and oxygen partial pressure that cause the passive to active transition at  the SiO2 -SiC  interface have been reported  in  the literature.41,47-51 When oxidation occurs  in air or under oxygen/inert mixture at 1600 C  the passive  to active  transition  is believed to happen if the oxygen partial pressure becomes lower −1 atm.47-49 In these particular environmental conditions than 10 several reactions of SiC active oxidation can occur: SiC(s) +  2SiO2(s) ↔  3SiO(g) + SiC(s) +  2SiO(g) + SiC(s) + SiC(s) + SiC(s) + SiC(s) +   SiO2(s) ↔  SiO2(s) ↔  O2(g) ↔  ½O2(g) ↔  ½O2(g) ↔   SiO(g) +  SiO(g) +   3Si(s,l) +   Si(s,l) +   2CO(g)   CO(g)  (3)  (4)  (5)  (6)  (7)  (8)   C(s)      CO(g)   C(s)   CO(s)  With the progressive increase of temperature and the progressive decrease of  the oxygen potential  the  transition should  initially involve the reaction (3).41,47 Therefore  in  this  case  the SiC oxidation process  can be described according to the Eq. (6). These reactions should cause the depletion of both Si and C owing to the formation of gaseous SiO and CO. The  results obtained  in  the present work  show that the transition involves the reaction (4) instead. This means that  the oxidation of SiC can be summarized by  reaction  (7). However also  impurities are believed  to have an effect on  the passive/active  transition.48 Dealing with  the materials under investigation  the presence  in  the  laminates of some percent of carbon, added as sintering aid or arising from thermal decomposition of binder and plasticizer, likely can affect the equilibrium reactions  reported above.  In particular  the  formation of carbon oxide coming from combustion of carbonaceous materials, which is additional to CO resulting from oxidation of SiC, could hinder reactions (3) and (6). Nevertheless the presence of carbon could also affect equilibrium reactions (4) and (7), but gaseous CO is expected to diffuse much more quickly than solid carbon and therefore affecting more the chemical equilibria. All the microanalysis techniques concurred to show that the porous layer arises from the passive oxidation of SiC to gaseous SiO and solid C, then the porous layer is a silicon depletion zone that still contains significant amount of carbon. This last finding seems  to be peculiar of  the  laminates under  investigation  that behave differently from other SiC-ZrB2 composites processed according  to other methods.  In  fact  in  the case of  these other similar composites  the porous  layer has been  referred  to  the depletion of both Si and C resulting from reaction (6) of active oxidation of SiC. The superior oxidation resistance of SiC and SiC-rich members of the SiC-ZrB2 system is expected not to exist any longer when the oxidation temperature exceeds about 1720 C. In fact     \\x0c', 'E. Padovano et al. / Journal of the European Ceramic Society 35 (2015) 1699-1714   1713     the passive  layer of  silica melts  and undergoes  evaporation when the temperature exceeds this limit. Over the silica melting point  the unique  solid oxidation product  for SiC-ZrB2 composites  is zirconia, which requires  the adoption of ZrB2 -based materials  for applications  in  these extreme conditions. These considerations drive  to  the concept of developing TPSs alternating layers based on SiC and ZrB2 respectively, since such a kind of architecture could well sustain oxidation under a variety of environmental conditions. In fact the consumption of a SiCbased barrier, occurring at  temperature over 1700 C or under active oxidation conditions, would result in the exposure to the aggressive atmosphere of a ZrB2 -rich layer that can stop or delay the TPS recession under these conditions. On the other hand, preliminary oxidation tests carried out on laminates with alternate layers of SiC and ZrB2 -based composite proved that the passive layer  formed on  the SiC  layer when  the  temperature  is  lower than 1600 C is able to preserve from oxidation the underlying sheets with high ZrB2 content. Laminates with composite layers alternating with those of SiC were successfully prepared, but the processing method should be improved in order to obtain robust laminates. To this purpose powders of SiC and ZrB2 with similar size distribution could be adopted and, in case, the content of binder  in  the composite slurry could be reduced  in order  to decrease the shrinkage.     5. Conclusions     A single fabrication method for processing several ceramics belonging to the SiC-ZrB2 system has been developed in view of applications for reusable TPSs of space vehicles. According to this method, exploiting tape casting and pressureless sintering, it was possible to process laminates of composition ranging from 100%SiC to 80 vol.%ZrB2-20 vol.%SiC as well as  laminates with alternating  layers of SiC and SiC-ZrB2 composites. Even  though  the suitability of  the method was assessed for these  last  laminates with alternate  layers,  their production process should be  further  improved  for avoiding  the presence of defects and enhancing the mechanical characteristics. The oxidation resistance in air up to 1600 C of the composite  laminates as well as  the oxide scale microstructure and  the residual strength and stiffness after oxidation depended on  the SiC:ZrB2 ratio. Actually the oxidation resistance progressively worsened with the ZrB2 content increase. Nevertheless  all  laminates  investigated  displayed  passive behavior, since the oxidation products obtained on the material surface were able  to slow down  the oxidation process. However cristobalite  layer granted better oxidation protection  than borosilicate glassy layers. SiC:ZrB2 volume ratios not lower than 80:20 resulted in the formation of a rather  thin cristobalite passive  layer on  the surface. Nevertheless impurities contained in the ceramic laminate favored  the growth of  the scale with respect  to what expected for high-purity SiC. The oxide scales grown on composite laminates with at least 40 vol.% of ZrB2 showed some differences from the oxide layers obtained by oxidizing similar monolithic composites, even     though in every case the scale consists of three sub-layers (external glassy  layer,  two-phase  intermediate  layer, porous  inner layer). In the case of laminates, surprisingly, the external glassy layer contained carbon inclusions and not negligible amount of boron even after oxidation in air at 1600 C for 24 h. The  intermediate  layer consisting of crystalline ZrO2 plus borosilicate glass showed the columnar structure, usually mentioned in the current models for ZrB2-SiC composite oxidation, only when the laminate contained 80%ZrB2 , but in all the other cases it consisted of two interpenetrated and randomly oriented networks of the two phases. A porous internal layer formed because of SiC active oxidation, that is SiC depletion, but this phenomenon occurs according to different reactions in laminates and corresponding monolithic composites.  In case of conventional ceramic composites both Si and C depletion was reported  in  literature, while  in case of ceramic laminates only Si depletion occurred and therefore the carbon percentage still  remained very high  inside  the porous layer. 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},{
  "_id": 150,
  "PDF": "Oxidation behavior of ZrB2-xSiC composites at 1500°C under different oxygen partial pressures.pdf",
  "Text": "['Available online at www.sciencedirect.com  Ceramics International 40 (2014) 15303-15311  CERAMICS  INTERNATIONAL  www.elsevier.com/locate/ceramint  Oxidation behavior of ZrB2-xSiC composites at 1500 1C under different oxygen partial pressures  Young-Hoon Seong, Do Kyung Kimn  Department of Materials Science and Engineering, KAIST, 291 Daehak-ro, Yuseong-gu, Daejeon 305-701, Republic of Korea  Received 16 May 2014;  received in revised form 7 July 2014; accepted 7 July 2014  Available online 14 July 2014  Abstract  The oxidation behavior of zirconium boride composites with various SiC contents (0-40 vol%) at 1500 1C in air (pO2 ¼ 104 Pa) and under low pO2 (10-8 Pa) was investigated. Due to different oxidation kinetics calculated from the oxidation depths, the oxidized composites exhibited different layered structures. In addition, the composites of ZrB2-30 vol% SiC (one of the typical compositions) oxidized at 1500 1C for 10 h in air and under low pO2 conditions were analyzed using TEM (transmission electron microscope). In air, the oxidation depth as a function of time indicated a parabolic kinetic behavior, and the ZrB2-40 vol% SiC composite exhibited the lowest parabolic rate constant (kP) of 232 μm2/h. Under low pO2, the oxidation depth as a function of time indicated a parabolic to linear transition kinetic behavior, except for monolithic ZrB2. The (kP) of 811 μm 2/h. monolithic ZrB2 exhibited the lowest parabolic rate constant & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  Keywords: ZrB2-SiC; Oxidation behavior; Oxygen partial pressure; Oxidation rate constant  1.  Introduction  The ultra-high temperature ceramics (UHTCs) exhibit several properties that are suitable for extreme environments including a high melting temperature (4 3000 1C), high mechanical properties, and chemical stability. Boride and carbide ceramics are among the UHTCs, and zirconium diboride (ZrB2) is advantageous for aerospace applications and thermal protection systems because ZrB2 has a low density (6.085 g/cm3) and high thermal conductivity (60-130 W/mK) [1-6]. SiC is widely used as an additive in ZrB2 to improve the oxidation resistance of ZrB2 due to the formation of a surface SiO2 layer, and the sintering behavior (sintering additive) and mechanical properties (crack deflection) are also improved [2,5-7]. Under high temperature oxidizing environments (at a 1500 1C), temperature of the four-layered structure (surface SiO2 region, oxide region, SiC-depleted region, matrix) are formed. The surface silica-rich layer protect the oxygen transport through the oxide region resulting in parabolic mass  nCorresponding authors. Tel.: þ 82 42 350 4118; E-mail address: dkkim@kaist.ac.kr (D.K. Kim).  fax: þ 82 42 350 3310.  http://dx.doi.org/10.1016/j.ceramint.2014.07.036  0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  gain kinetics [1,8]. The surface SiO2 region is dense and continuous silica-rich phase. The oxide region consists of ZrO2 and amorphous SiO2 oxide, and SiC-depleted region consists of ZrB2 and pores. These three regions (surface SiO2 region, oxide region, SiC-depleted region) are referred to oxide scales. Rezaie et al. reported the pO2 boundary conditions of the SiC oxidation at 1500 oC using the thermodynamic calculations is below 8.8 \\x02 10 \\x00 13 Pa, (volatility diagram). When the pO2 the SiC will be removed by the phase transformation from SiC 8.8 \\x02 10 (s) to SiO (g). On the contrary, when the pO2 is above \\x00 13 Pa, the SiC will oxidized to SiO2 (l) which will then volatilize to SiO (g) [9]. The pores originated from the evaporation of SiC (s) were created, and this ZrB2 region containing pores is called ‘SiC-depleted region’. Under re-entry conditions, the pO2 level decreases because O2 collides with the hypersonic flight vehicle, and O2 is dissociated to form O atoms [9]. The change in the pO2 level is one of the most important parameters for high temperature oxidation of ZrB2-SiC composites. Li et al. studied the pO2 effect on the high temperature oxidation of the ZrB2-SiC composite. A low pO2 promotes the active oxidation of SiC, and the oxidized products and microstructure of the oxidized  \\x0c', '15304  Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  region under O conditions are similar to those in O2 conditions. Due to the oxygen chemical potential gradient, the oxidation behavior or kinetics of SiC is highly sensitive to the pO2 in the reaction layer [10]. Han et al. studied the oxidation behavior of ZrB2-SiC at 1800 1C under different pO2 (0.2 and 2 \\x02 10 \\x00 4 atm). They pointed out that low oxygen partial pressure was detrimental to the oxidation resistance of the composites because of the active oxidation of SiC. The low pO2 decreased the high temperature oxidation resistance of the ZrB2-SiC composites due to the active oxidation of SiC [11]. In this work, we investigated the effects of SiC addition and oxygen partial pressure on high temperature oxidation behavior of ZrB2 with various SiC contents at 1500 1C under different pO2 by high resolution microstructure analysis (TEM). Also, the high temperature oxidation kinetics of ZrB2 with various SiC contents under different pO2 was studied by oxidation rate constants calculation from oxidation depth.  2. Experimental procedure  2.1. Preparation  (Hexagonal, 4 99%, 3-5 μm, Grade-A, Commercial ZrB2 and SiC (α-phase, Hexagonal, H.C. Starck, Germany) 6 H0.45 μm, UF-25, H.C. polytype, 98.5%, Starck, Germany) were used. Four batches consisted of 60, 70, 80, and 100 vol% ZrB2 powder with 40, 30, 20, and 0 vol% SiC powder, respectively, and they were abbreviated as Z0S, Z2S, Z3S, and Z4S in the following discussion. ZrB2 was mixed with SiC and wet milled for 24 h with ZrO2 balls. The milled powders were heated with stirring and granulated using a 325 mesh sieve. The mixed powders were sintered in a hot press (Thermal Technology Inc., HP20-1000-3560, CA, USA) at 1950 1C with a uniaxial pressure of 32 MPa for 2 hours under specimens were diced (4 \\x02 4 \\x02 3 an argon atmosphere. The mm) to prepare samples for the high temperature oxidation test, and all of the information regarding the specimens is listed in Table 1.  2.2. Oxidation  A high temperature oxidation test was conducted on a tube furnace equipped with MoSi2 heating elements. Prior to the tests, the sintered samples were polished and inserted in the middle of the furnace with an alumina boat. The high temperature oxidation test was performed at 1500 1C for 1, 3,  5, 7, and 10 h in air and under low pO2. The CO/CO2 mixed gas (CO2: 0.002 mol) was used to produce a low pO2 \\x00 8 Pa) at 1500 1C [9]. Based on the information obtained (10 \\x00 8 Pa was from a thermodynamic model, a pO2 value of 10 selected to form a SiC depletion region when the ZrB2-SiC is oxidized [9,12]. For the mixed gas flow, an Al2O3 tube furnace was employed with a gas flow rate of 50 cm3/min.  2.3. Characterization  According to the Archimedes principle, the bulk densities of the consolidated billets were calculated. To determine the theoretical densities of each composite, we used the rule of mixtures and the density values of the pure phases (i.e., ZrB2: 6.09 g/cm3 3.21 g/cm3). and SiC: FE-SEM (Philips, XL30 FEG, Eindhoven, Netherlands) was used to analyze the microstructures of vertical sections of the oxidized composites. The specimens used in the TEM observations were prepared using conventional mechanical polishing and an ion polishing system (Focused Ion Beam, Quanta 3D FEG, FEI, Eindhoven, Netherlands). TEM (Tecnai G2 F30 S-twin, 300 kV, FEI, Eindhoven, Netherlands) was used to acquire the bright field (BF) images, selected area electron diffraction (SAED) patterns and elements maps.  3. Results and discussion  3.1. Density and microstructure  Based on the rule of mixtures principle, the calculated theoretical densities of the composites were 6.09, 5.51, 5.23, and 4.94 g/cm3 for Z0S, Z2S, Z3S, and Z4S, respectively. The relative densities of the composites were 98.7, 99.1, 99.7, and 97.4% for Z0S, Z2S, Z3S, and Z4S, respectively. All of the composites have relative densities over 97%, which indicates that the porosity will not significantly affect the high temperature oxidation kinetics. Fig. 1(a) shows a TEM image (bright field) of the as-sintered ZrB2-30 vol% SiC specimen. The dark gray grains are ZrB2 particles, and the light gray grains are SiC particles. The ZrB2 and SiC particles have a faceted shape with a grain size of 1.0-3.0 μm. The SAED (Selected Area Electron Diffraction) patterns obtained from the circled areas are shown in Fig. 1(b). The [0111 ] zone axis pattern of hexagonal ZrB2 and the [02 21] zone axis pattern of SiC-4H (hexagonal) are shown in Fig. 1(b). The SiC-4H was partially transformed from the SiC-6H phase during sintering because  Table 1  Summary of ZrB2 x vol% SiC composite specimens: Compositions, Designations, Relative Density, and Conditions of Oxidation Tests.  Designations  Composition (ZrB2:SiC, volume ratio) Relative density (%)  Oxidation test conditions  Oxygen partial pressure (Pa)  Temperature (1C)  Time (h)  high pO2 low pO2  Z0S  100:0  Z2S  80:20  99.1  98.7 2 \\x02 104 (air) \\x00 8 (CO/CO2 mixed gas) 10 1500  1-10  Z3S  70:30  99.7  Z4S  60:40  97.4  \\x0c', 'Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  15305  Fig. 1. The TEM (a) BF image and (b) SAED patterns of as-sintered ZrB230 vol% SiC. The SAED patterns were taken from circled areas. These patterns show the ½011 1 \\x8a zone axis pattern of ZrB2 and the ½02 21\\x8a zone axis pattern of SiC.  stable polytypes at  the SiC-4H and -6H phases are the most 1800-2000 1C [13]. The SEM images in Fig. 2 show the hot-pressed ZrB2-SiC composites. The pores were rarely observed. In addition, the light gray grains in the micrographs are ZrB2, and the dark gray grains are SiC. In general, SiC was well dispersed in the ZrB2 composites containing SiC even though the SiC was locally aggregated.  3.2. Oxidation in high pO2  The cross-sectional SEM images in Fig. 3 show the oxidized (pO2 ¼ composites tested at 1500 1C for 10 h in air 2 \\x02 104 Pa). To compare the microstructures and oxidation depths, the magnifications shown in Figs. 3 and 4 were held constant (except Z0S). The reaction region (surface amorphous region þ oxide silicon oxide region) depths of ZrB2-x vol% SiC (x=0, 20, 30 and 40) were 1150, 65.5, 52.2 and 47.6 μm, respectively. The oxidation of ZrB2-x vol% SiC at 1500 1C produced a three layered structure composed of a surface amorphous silicon oxide region, an oxide interlayer region, and an unreacted region except for monolithic ZrB2. The reaction depths of the SiC containing composites decreased slightly as the SiC content increased. In case of monolithic ZrB2, no oxygen diffusion barrier exists on the surface or at the ZrO2 grain boundaries due to the absence of a Si source that could form a silica rich phase. Therefore, a drastic increase in the reaction depth is affected by large numbers of oxygen paths (porous ZrO2 structure, inset of Fig. 3(a)). Some previous studies have reported that the four-layered structure (i.e., surface SiO2, oxide region, SiC-depleted region, matrix)  was created by Reactions (1) and (2) after oxidation tests at 1500 1C [1,7-9,14]. However, in this study, the SiC-depleted region was not observed, and a different three-layered structure (i.e., surface SiO2, oxide region, matrix) was formed. The pO2 level at the interface between the oxide region and the unreacted region cannot be determined exactly but it is above 8.8 \\x02 10 \\x00 13 Pa. This value is a boundary condition for Reactions (2) and (3) based on the pO2 for the SiC-SiO and SiC-SiO2 equilibrium at 1500 1C [9,12], and the SiC depleted region can be formed below 8.8 \\x02 10 \\x00 13 Pa by Reaction (3). The SiC-depleted region in the ZrB2-SiC composite is generated from SiC dispersion in the composite [15]. In certain cases, SiC-depleted region was not created due to the volume ratio of SiC and ZrB2. The high volume fraction of the SiC grains leads to high interconnectivity of the SiC grains, and the number of isolated SiC grains decreased (Fig. 1(a) and Fig. 2 (c), (d)). Therefore, SiC was not converted to SiO (g) but oxidized to form an amorphous SiO2 at 1500 1C that flowed viscously along the ZrO2 grain boundaries. The surface SiO2 scale could protect the matrix from inward transport of oxygen during oxidation because SiO2 is less volatile than B2O3 [1,8,9]. The volume expansion due to oxidation of SiC might be one of the driving forces for viscous flow of SiO2 to the surface [16]. The surface SiO2 oxide region in Z4S was more easily formed than in the other cases, and this region effectively prevented inward oxygen transport in the composite. ZrB2 (s) þ 5/2O2 (g)-ZrO2 (s) þ B2O3 (g)  (1)  SiC (s) þ 3/2 O2 (g)-SiO2 (l) þ CO (g)  (2)  \\x0c', '15306  Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  Fig. 2. SEM images of  the as-sintered ZrB2 x vol% SiC composites (a) x ¼ 0,  (b) x ¼ 20,  (c) x ¼ 30, and (d) x ¼ 40.  Fig. 3. SEM images of  the ZrB2 x vol% SiC composites oxidized at 1500 1C for 10 h in air  (a) x ¼ 0,  (b) x ¼ 20,  (c) x ¼ 30, and (d) x ¼ 40.  \\x0c', 'Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  15307  Fig. 4. SEM images of  the ZrB2 x vol% SiC composites oxidized at 1500 1C for 10 h under  low pO2 (a) x ¼ 0,  (b) x ¼ 20,  (c) x ¼ 30, and (d) x ¼ 40.  SiC (s) þ O2 (g)-SiO (g) þ CO (g) SiO2 (l)-SiO (g) þ 1/2 O2 (g) SiC (s) þ 1/2O2 (g)-SiO (g) þ C (s)  (3)  (4)  (5)  3.3. Oxidation in low pO2  Under low pO2, in contrast to the results in air, the surface of the composite was not covered with a silica-rich scale, and two regions (oxide region, matrix) were observed. The crosssectional SEM images of the oxidized composites tested at 1500 1C for \\x00 8 Pa) Fig. 4. The reaction region depths of ZrB2-x vol% SiC (x ¼ 0, 10 h under low pO2 (10 are shown in 20, 30 and 40) were 93.1, 122, 147 and 172 μm, respectively. Rezaie explained that the low pO2 promotes active oxidation of SiC due to the formation of SiO (g) and CO (g), and the active oxidation of SiC produced a SiC-depleted region [9]. Under low pO2, the SiC directly evaporated (SiC (s) SiO \\x00 15 Pa o pO2 o 8.8 \\x02 10 \\x00 13 Pa) or (g), Reaction (3), 10 evaporated in stages (SiC (s) SiO2 (l) SiO (g), Reaction (1) and (4), 8.8 \\x02 10 \\x00 13 Pa o pO2o 10 \\x00 5 Pa). In addition, ZrB2 stable below pO2 \\x18 1.9 \\x02 10 \\x00 11 Pa or oxidized to form was ZrO2 and B2O3. Therefore, the high temperature oxidation of the ZrB2-SiC composites was sensitive to the pO2 level. For Z4S, the oxygen was more easily transported inside the matrix because it has a higher porosity due to SiO vaporization from the grain boundary of the SiO2 phase compared to the other cases.  3.4. TEM analysis of oxidized ZrB2-xSiC  The TEM BF images in Fig. 5 show the microstructure of the oxide region of the Z3S composite oxidized in air and under low pO2. Fig. 5(a) and (b) were captured from the oxide region (mid-layer) just above the unreacted matrix region. Although the TEM specimens were taken from the interface between the unreacted layer and oxidized layer, the SiC depletion region was not found in both cases (air and low pO2). The microstructures of the oxide region were similar, though the pO2 values in the oxidation condition were different. Nearly all of the ZrB2 grains were oxidized and separated into small grains where their shape changed from faceted to round. As a result of grain dividing, many grain boundaries were created, which may play a role of the oxygen transport path. However, the SiC grains were partially oxidized and transformed to an amorphous SiO2 phase, and unoxidized SiC remained in the middle of the amorphous SiO2. By the ZrB2-ZrO2 equilibrium at 1500 1C analyzing the pO2 for from the volatility diagram [9,17], the pO2 in the oxide region 8.8 \\x02 10 \\x00 13 Pa. is thought to be above In Fig. 5(b), the amorphous SiO2 was dispersed along the ZrO2 grain boundary \\x00 8 Pa). in the Z3S composite tested under low pO2 (10 Although the overall microstructure was similar to that in the oxide region for Z3S tested in air (Fig. 5(a)), pores were observed at the grain boundaries. The SiO (g) evaporated from amorphous SiO2 leaving some pores at the grain boundary because the pO2 of this region was sufficiently low. From the thermodynamic calculations, the phase transformation from SiC (s) to amorphous SiO2 (l) (Reaction 2) is followed by SiO  \\x0c', '15308  Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  Fig. 5. TEM BF images of the oxide layer (just above the unreacted layer) of the ZrB2-30 vol% SiC composite oxidized at 1500 1C for 10 h (a) in air and (b) under low pO2. The STEM images and their elemental maps of the oxidized ZrB2 -30 vol% SiC composites (a) in air and (b) under low pO2.  (l) at a pO2 of 8.8 \\x02 10 \\x00 13 Pa to (g) evaporation from SiO2 \\x00 5 Pa (Reaction 4) 10 [9]. The pores in the amorphous SiO2 originated from the evaporation of SiO (g), and pSiO was calculated to be 38.2 Pa using the pO2 of the test condition \\x00 8 Pa) (10 in Reaction (4) [12]. Therefore, the oxidation 1500 1C behavior and kinetics of the ZrB2-SiC system at under low pO2 are much different from those of the ZrB2-SiC system in air. The elemental analyses in the oxide layer of the Z3S composite oxidized in air and under low pO2 including the partially oxidized SiC grain are shown Figs. 5(c) and 5(d), respectively. For the Z3S composite oxidized in air, Si and O were detected around the unreacted SiC while Si and C were detected in the unreacted SiC grain. Based on the SEM results, the Z3S composite that was oxidized in air did not create a SiC-depleted layer, which indicates that the O2 transport in this  area was relatively unchanged and the pO2 in the oxide layer as not low enough to allow for active oxidation (Reactions 3- 5). Namely, passive oxidation (reaction 2) was the predominant reaction in this layer. In contrast, for the Z3S composite oxidized under low pO2, C was detected with Si and O around the unreacted SiC. As the pO2 decreased, the active oxidation of SiC was classified as direct active oxidation (Reaction 3, below pO2 \\x18 8.8 \\x02 10 \\x00 13 Pa) and indirect active oxidation (Reactions 2 and 4, above pO2 \\x18 8.8 \\x02 10 \\x00 13 Pa) [9]. Based \\x00 8 Pa) on the pO2 (10 level of this experimental condition, indirect active oxidation (Reactions 2 and 4) most likely occurred first, and then, the C was dissociated from SiC (Reaction 5) due to a lower pO2 from the additional consumption of O2 during indirect active oxidation of SiC (Reactions 2 and 4) and oxidation of ZrB2 (Reaction 1). Based on the detection of carbon and the thermodynamic calculations, the  \\x0c', 'Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  15309  in this area may reach \\x18 5 \\x02 10 \\x00 15 Pa [12]. These pO2 level results are consistent with those obtained by Ni [18], Carney [19] and Monteverde [20].  3.5. Oxidation kinetics of  the ZrB2-xSiC  Fig. 6 shows the oxidation kinetics of the ZrB2-x vol% SiC specimens oxidized in an air atmosphere at 1500 1C. The reaction depths as a function of oxidation time are shown in Fig. 6(a). In air, the plots indicate that oxidation of the ZrB2SiC composites follow a parabolic rate law. The silica-rich scales located on the surface and grain boundaries are dense, smooth and viscous, as shown in Fig. 3. Therefore, the silica rich scale is protective, and the oxidation kinetics are controlled by oxygen diffusion. The overall linear shape of the squared depth increase as a function of time in Fig. 6(b) confirms that the oxidation kinetics follow a parabolic rate law. The parabolic rate (kP) calculated according to d2 ¼ kPt constants (t: oxidation  time) are 1.25 \\x02 105, 433, 269 and 232 μm2/h for Z0S, Z2S, Z3S and Z4S, respectively [21]. The ZrB2-40 vol% SiC (Z4S) composite exhibited the lowest parabolic rate constant (kP). In air, the composite with a higher SiC content exhibits enhanced oxidation resistance at 1500 1C. The oxidation kinetics of the ZrB2-x vol% SiC specimens oxidized under low pO2 at 1500 1C are shown in Fig. 7. The reaction depths as a function of time for the studied specimens are shown in Fig. 7(a). The oxidation depths at 1500 1C as a function of time exhibit a parabolic-linear transition shape. During the initial 3 h, parabolic oxidation kinetics were observed followed by linear oxidation kinetics, except in the case of Z0S. The time when the slope (derivative) of the parabolic function is equal to the slope of the linear function was determined as the parabolic-linear transition time. In this study, the parabolic-linear transition time is 2.9768, 3.0327 and 2.9846 h after reaching the maximum temperature (1500 1C) for Z2S, Z3S and Z4S, respectively. This result  Fig. 6. Oxidation behavior of  exposed to air  for  the ZrB2 x vol% SiC composites at 1500 1C a period of time up to 10 h. (a) Parabolic plots for the  Fig. 7. Oxidation behavior of ZrB2 composites with various amount of SiC (in vol%) at 1500 1C for a period of time up to 10 h under a low oxygen partial  pressure.  (a) Parabolic-linear  transition plots  (except Z0S)  for  the oxidation  oxidation depth as a function of  time and (b)  linear plots  for  the square of  depth as  a  function of  time and (b)  linear-quadratic  transition plots  (except  oxidation depth as a function of  time.  Z0S)  for  the square of oxidation depth as a function of  time.  \\x0c', '15310  Table 2  Calculated parabolic and linear rate constants of the ZrB2 composites at 1500 1C under different oxygen partial pressures.  x  vol% SiC  Composites Atmosphere  Air  Low pO2  kP  (μm2/h)  Error  limits  kP (μm2/h, 0-3 h)  Error  limits  kL (μm/ h,  3-10 h)  Error  limits  Z0S  Z2S  Z3S  Z4S  125,442  12,610  811 (1-  24.61    433.25  269.43  231.61  11.74  14.85  6.40  10 h)  1,007  1,824  2,206  5.81  8.78  181.61  9.262  10.972  11.971  0.2510  0.1478  0.1202  indicated that the dominant reaction changed from oxidation due to oxygen diffusion (Reactions 1 and 2,4 1300 1C) to range of 8.8 \\x02 dissociation of SiO from SiO2 (Reaction 4; \\x00 13 o pO2 o 10 \\x00 5 Pa 10 at 1500 1C) [9,16]. The paraboliclinear transition behavior can be explained by an increase in defect (connected pores and cracks) interconnectivity due to SiO dissociation from amorphous SiO2 (Fig. 5(b)) [22]. The equilibrium vapor pressure of SiO (g) at 1500 1C was relatively high ( \\x18 38.2 Pa). When the defects interconnect with each other to form oxygen transport paths, the parabolic-linear transition of the oxidation kinetics occurred. The linear-quadric shape of the squared depth increase as a function of time in Fig. 7(b) confirms that the oxidation follows the parabolic rate law-linear rate law transition after 3 h of oxidation. The calculated parabolic rate constants 1.01 \\x02 103, 1.82 \\x02 103 (kP, 0-3 h) are 811 (1-10 h), and 2.21 \\x02 103 μm2/h for Z0S, Z2S, Z3S and Z4S, respectively. The linear rate constants (kL, 3-10 h) calculated according to 11.97 μm/h the relationship are 9.26, 10.97 and for Z2S, Z3S and Z4S, respectively. The monolithic ZrB2 (Z0S) exhibited the lowest parabolic rate constant (kP) of 811 μm2/h. The composite with a lower SiC content exhibited enhanced oxidation resistance at 1500 1C under low pO2 conditions. The oxidation rate constants and their error limits of all of the composites at 1500 1C are summarized in Table 2.  d ¼ kLt  4. Conclusions  In the present work, the high temperature oxidation behavior and kinetics of ZrB2 with various content of SiC (ZrB2-xSiC) were studied by TEM analysis and oxidation rate constants calculation. These results will help to the better understanding of the high temperature oxidation behavior and mechanism of the ZrB2 composites containing SiC. The hot-pressed ZrB2-x (pO2 ¼ 2 \\x02 104 Pa) vol% SiC composites were oxidized in air \\x00 8 Pa) at 1500 1C. After oxidation, a and under low pO2 (10 three-layered structure (i.e., surface SiO2, oxide region, and matrix) was observed in the SiC containing composites oxidized in air. However, two regions (i.e., oxide region, matrix) were observed in all of the composites oxidized under low pO2 with varying depths. The SiC-depleted region was not observed in both cases (air and low pO2). The inhomogeneous  Y.-H. Seong, D.K. Kim / Ceramics International 40 (2014) 15303-15311  dispersion of the SiC in ZrB2, high volume fraction of the SiC (the number of interconnected SiC grains increased, the number of isolated SiC grains decreased) and the internal pO2 level (not low enough to form the SiC-depleted region) could be the possible reasons. Based on the TEM results, the amorphous SiO2 and unreacted SiC were dispersed in the oxide region of ZrB230 vol% SiC. The faceted ZrB2 grains were separated into small spherical ZrO2 grains after oxidation. The unreacted SiC existed in the amorphous SiO2 as an island structure. Although the overall microstructure was similar in both cases, pores were observed at the grain boundaries in the Z3S composite tested under low pO2 due to evaporation of SiO (g) from amorphous SiO2. From the EDS observations, the predominant reaction (i.e., oxidation) is highly dependent on the oxygen partial pressure. In air, the oxidation depth as a function of time exhibited a parabolic kinetics behavior, and the oxidation rate constants were calculated from the oxidation depth results. The ZrB2- 40 vol% SiC composite exhibited the lowest parabolic rate constant (kP) of 232 μm2/h. Under low oxygen partial pressure, the oxidation depth as a function of time exhibited a parabolic to linear transition kinetics behavior except in the case of monolithic ZrB2. The monolithic ZrB2 exhibited the lowest parabolic rate constant (kP) of 811 μm 2/h. The oxidation resistance of the ZrB2-x vol% SiC composites at 1500 1C is highly dependent on the SiC content and oxygen partial pressure. The high temperature oxidation behavior and kinetics of ZrB2 with various content of SiC (ZrB2-xSiC) were studied by TEM analysis and oxidation rate constants calculation, respectively.  Acknowledgements  We would like to acknowledge the financial support from R&D Convergence Program of MSIP (Ministry of Science, ICT and Future Planning) and ISTK (Korea Research Council for Industrial Science and Technology) of Republic of Korea (Grant B551179-13-03-01). Also, this work was supported by the Defense Acquisition Program Administration and Agency for Defense Development under contract UD140023GD, and the Priority Research Centers Program through the NRF funded by MEST (2012-048034).  References  [1] W.G. Fahrenholtz, G.E. Hilmas,  I.G. Talmy,  J.A. Zaykoski, Refractory  diborides  of  zirconium and  hafnium,  J. Eur. Ceram. Soc.  90  (2007)  1347-1364.  [2] M.J. Gasch, D.T. Ellerby, S.M. Johnson, Ultra high temperature ceramic  composites,  in: N.P. Bansal  (Ed.), Handbook of Ceramic Composites,  Springer, New York, 2004, pp. 197-224.  [3] M.M. Opeka, I.G. Talmy, J.A. 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Monteverde, Beneficial effects of an ultra-fine α-SiC incorporation on the sinterability and mechanical properties of ZrB2, App. Phys. A 82 (2006) 329-337.  [21] L.F. He, Z.J. Lin, Y.W. Bao, M.S. Li, J.Y. Wang, Y.C. Zhou, Isothermal oxidation of bulk Zr2Al3C4 at 500 to 1000 1C in air, (2008) 359-366.  J. Mater. Res. 23  [22] D. Jayaseelan, E. Zapata-Solvas, P. Brown, W.E. Lee,  In situ formation  of oxidation resistant  refractory coatings on SiC-reinforced ZrB2 ultra high temperature ceramics, J. Am. Ceram. Soc. 95 (2012) 1247-1254.  \\x0c']"
},{
  "_id": 151,
  "PDF": "Oxidation behavior of ZrB2–MoSi2–SiC composites in air at 1500 C.pdf",
  "Text": "['Available online at www.sciencedirect.com  Ceramics International 37 (2011) 585-591  www.elsevier.com/locate/ceramint  Oxidation behavior of ZrB2-MoSi2-SiC composites in air at 1500 8C  Shuqi Guo a,*, Takashi Mizuguchi a, Masahide Ikegami b, Yutaka Kagawa a,b  a Hybrid Materials Center, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan  b Research Center for Advanced Science and Technology, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan  Received 2 July 2010; received in revised form 20 August 2010; accepted 26 September 2010  Available online 28 October 2010  Abstract  We investigated the oxidation behavior and the effect of the amount of SiC added on oxidation resistance in both hot-pressed ZrB2-MoSi2-SiC composites, 55ZrB2-40MoSi2-5SiC and 40ZrB2-40MoSi2-20SiC (vol.%), exposed to dry air at 1500 8C for up to 10 h. Quantitative electron  microprobe analysis characterizations of the chemical compounds of post-oxidized composites were carried out. Parabolic oxidation behavior was  observed for both composites. The addition of SiC improved the oxidation resistance of ZrB2-MoSi2-SiC composites, and the improvement  enhanced with amount of SiC added. The microstructure of the post-oxidized composites consisted of two characteristic regions: oxidized reactive  region and unreactive bulk material  region. The oxidized reactive region divided into an outermost dense silica-rich scale layer and oxidized  reactive mixture layer. The improvement of oxidation resistance with SiC addition is associated with the presence of a thicker dense outermost  scale layer which inhibited inward diffusion of oxygen through it. # 2010 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  Keywords: Ceramics; Corrosion; High-temperature properties; Electron microprobe analysis  1.  Introduction  resulting oxidized reactive region consists of  three layers:  (i)  the outermost glassy layer, (ii) the oxide subsurface layer, and  Zirconium diborides (ZrB2)-based composites have become  (iii) the SiC-depleted ZrB2 layer. For the ZrB2-MoSi2 system,  an important class of materials  for  structural applications at  on the other hand, MoSi2 is oxidized in atmospheric air, instead  ultra high temperatures because they have an extremely high melting point (>3000 8C) [1-4]. A major problem with ZrB2 based ceramics is high temperature oxidation where they are  pressurelessly sintered 700-1400 8C in  dry  of SiC. Sciti et al. [11,12] examined the oxidation behavior of  20 vol.% MoSi2-containing ZrB2  at  air. They  showed  that  the  oxidation  considered to be applied as structural materials for use in high resistance was  improved with MoSi2  added,  as  a  result of  temperature oxidizing environments [5-7]. Heating ZrB2 in air  formation of SiO2 scale layer. The oxidation products consist of  produces a scale composed of ZrO2 and B2O3 in which B2O3 has a high vapor pressure and is vaporized above 1300 8C [5-7].  SiO2, ZrO2, ZrSiO4, MoO2, and MoB, and the oxidized reactive  region consists of a SiO2-rich glass layer; a subsurface oxide  Thus, the oxidation resistance of ZrB2 must be improved for use in oxidizing environments above 1300 8C.  layer; and a MoB, ZrO2, and SiO2-containing mixture layer,  depending  on  the  exposure  temperature.  These  studies  To improve oxidation resistance, Si-containing additives,  such as SiC and MoSi2, are added to ZrB2 [7-12],  forming a  demonstrated that  the oxidized reactions of ZrB2 depended  on the composition, exposure temperature, and exposure time at  protective borosilicate 1200 8C that enhances  glass  layer  at  temperatures  above  a particular temperature, and oxygen content  in the oxidizing  the oxidation resistance of ZrB2. For  atmosphere.  the ZrB2-SiC system, the main oxidation products are ZrO2, B2O3, and SiO2 below 1300 8C, ZrO2 and SiO2 above 1300 8C,  A recent study has shown that bending strength and fracture  toughness are better for ZrB2-MoSi2-SiC system than for both  as B2O3 liquid gets completely vaporized [5-10]. Generally, the  ZrB2-MoSi2 system and ZrB2-SiC system [13]. However, there  * Corresponding author. Tel.: +81 029 859 2223; fax: +81 029 859 2401.  E-mail address: GUO.Shuqi@nims.go.jp (S. Guo).  is a little known of the oxidation behavior of ZrB2-MoSi2-SiC  system. It could be expected that the oxidation behavior of the  ZrB2-MoSi2-SiC system differs with both the ZrB2-MoSi2 and  ZrB2-SiC systems. In this study, two hot-pressed compacts of 5  0272-8842/$36.00 # 2010 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  doi:10.1016/j.ceramint.2010.09.034  \\x0c', 'and 20 vol.% SiC-modified 40 vol.% MoSi2-containing ZrB2 were exposed in dry air at 1500 8C for up to 10 h. Quantitative  characterization of  the oxidized composites were conducted  using electron probe micro analysis  (EPMA)  to identify the  oxidation products, reactive compound composition, and their  distribution. Also,  the  effect  of  amount  of  SiC added  on  oxidation behavior was examined.  2. Experimental  2.1. Oxidation test  The material used in this study was prepared by hot-pressing  SiC-containing 40 vol.% MoSi2-ZrB2 composites. In order to  examine the oxidation resistance of ZrB2-MoSi2-SiC compo sites as well as to learn the effect of amount of SiC added, two  compositions of 5 and 20 vol.% SiC-modified 40 vol.% MoSi2- ZrB2 composites were hot-pressed at 1800 8C and 30 MPa in  vacuum for 30 min. The detailed sintering process has been  reported  elsewhere  [14]. Hereafter,  the  two  compositions  materials are denoted as ZMS5 and ZMS20.  Test specimens 2.5 mm \\x02 2 mm  with  average  dimensions  of  5 mm \\x02  were  cut  from  the  hot-pressed  ZrB2-  MoSi2-SiC composites plates with a diamond wafered blade.  After the specimens were polished with a diamond paste down  to 1.0 mm, they were ultrasonically cleaned in acetone and then temperature of 100 8C prior  kept  in an oven at a constant  to  oxidation. Oxidation tests were performed using an electronic  furnace (BFT-150-P, Nikkato Co., Ltd., Tokyo, Japan) at 1500 8C for up to 10 h in dry air. The heating and cooling rates 20 8C and 10 8C/min,  were  respectively. Before  and  after  oxidation,  the specimens were weighed, respectively, using an  analytical  balance  (AB265-S, Mettler  Toledo  Co.,  Ltd.,  Switzerland) with an accuracy of 0.1 mg.  2.2. Characterization  X-ray  diffraction  (XRD)  was  used  to  identify major  crystalline  phases  present  in  both  the  as-received  and  the  post-oxidized composites. The oxidized surfaces were  then  characterized  by  scanning  electron microscopy  (SEM)  and  energy-dispersive X-ray  spectroscopy  (EDX).  In  order  to  examine the oxidation evolution, the oxidized specimens were  cut  in half, and one of  the parts was mounted in epoxy and  carefully  polished with  a  diamond  paste  down  to  1.0 mm.  Cross-sectional observations of  the oxidized sample’s micro structures were  conducted  using  a  prototype wavelength  dispersive-electron probe microanalyzer (WDS-EPMA) which  was developed by Kimura et al. [15] based on a commercially  available FE-EPMA (JEOL;  JXA8900R).  In  addition, map  analysis of O, C, Si, Mo, Zr, and B elements within the cross section plane was conducted at 0.2 mm/step in the pixels of 200 \\x02 1024 using an X-ray mapping technique with EPMA,  with an accelerating voltage of 10 kV and probe current of 5.0 \\x02 10 \\x008 A. The X-ray image data obtained is projected  on  a  two-dimensional  space  of X-ray  intensity  to  form a  histogram, which is generally called a scatter diagram. Detailed  calculations  of  the  scatter  diagram have  been  reported  elsewhere [16]. The analysis of  the scatter diagram,  i.e.,  the  scatter diagram method, was used to identify the compounds  composition as well as determine their distribution, with the  pure SiO2, ZrO2, SiC, ZrB2 and MoSi2 powders as reference  materials.  3. Results and discussion  3.1. Weight gain  In Fig. 1, the plots of weight gain as a function of exposure  time for  the composites are presented. The two compositions  show similar oxidation behaviors: the specific weight increases  rapidly within  initial  1 h  of  exposure  and  then  the weight  increases gradually with exposure time, independent of the SiC  amount. This suggests that the oxidation mechanism is the same  for the two compositions. However, the weight gain was higher  for ZMS5  than  for ZMS20,  showing  the  improvement  of  oxidation resistance with SiC added. Their gains are approximately 3 and 6 mg/cm2,  specific weight  respectively, after  10 h of oxidation exposure. These gains are lower  than those  obtained in the 10-40 vol.% MoSi2-containing ZrB2 compo sites  exposed  at  the  same  oxidation  conditions which was  previously reported by authors elsewhere [17]. Fig. 2 presents plots of the square of the weight gain, W2, as a  function of oxidation time, t, for the composites. The oxidation  of the studied materials does not display parabolic kinetics at  the initial stage of oxidation (within 1 h). The deviation from  parabolic behavior suggests that  the oxidation behavior of the  studied materials is not appeared a parabolic behavior at  the  initial  stage. After oxidation of 1 h, however,  the parabolic  behavior is observed for each curve. This suggests that after the  outermost glassy scale formed at the initial stage of oxidation,  the diffusion, such as outward diffusion of constituent elements  from the bulk material  to the oxidized region and the inward  diffusion  of  oxygen  through  the  scale  layer,  is  the  rate controlling mechanism for  oxidation,  as  observed  for  both  ZrB2-MoSi2  and  ZrB2-SiC composites  [9,10,12].  Similar  oxidation behavior was previously reported in the 10-40 vol.%  10  2  )  1500°C  ZMS5  ZMS20  8  4  6  W  e  i  h g  t  G  a  i  n  ,  W  (  m  g  /  c  m  2  Time at Temperature, t (h)   0  12  10  8  6  4  2  0  Fig. 1. Plots of weight gain as a function of exposure time for the ZrB2-MoSi2- SiC composites oxidized at 1500 8C.  S. Guo et al. / Ceramics International 37 (2011) 585-591  586  [ ( ) T D $ F I G ]       \\x0c', 'MoSi2-containing  ZrB2  composites  oxidized  at  the  same  oxidation condition [17]. To calculate parabolic rate constant  of  the  studied materials,  it  is  assumed that  these materials  follow the parabolic behavior  for  the period of 1-10 h. The  parabolic  oxidation  rate  constants  of  two compositions, lines, are \\x183.6 and \\x181.1 mg2 cm obtained from the slopes of \\x004 h \\x001 for ZMS5 and ZMS20, respectively. These the straight \\x004 h \\x001) (4.4 mg2 cm  values  are  lower  than that  for SiC-free  40 vol.% MoSi2-containing ZrB2 composites oxidized at  the  same oxidation conditions [17]. Thus,  the added SiC to ZrB2-  MoSi2 system further improved its oxidation resistance, and the  improvement is enhanced with increasing amount of SiC added.  3.2. Microstructure of oxidized samples  Fig. 3 shows XRD patterns of the two compositions before  and after oxidation exposure. Before oxidation, only ZrB2,  MoSi2  and  SiC are  detected  for  each  composition. After  oxidation, a new primary oxidized phase of ZrO2 is detected in  the two compositions. The minor oxidized phases of MoB and  ZrSiO4 phases were present as well.  In addition, a trace of  amount of SiO2 phase was shown only in ZMS20. The intensity  of ZrO2 phase peaks decreases with the amount of SiC added,  showing improved oxidation resistance, because of  the ZrO2  phase resulting from oxidation of ZrB2. This agrees with the  decrease in weight gain observed with increasing amount of  SiC added (Fig. 1). For comparison, the intensity of ZrB2 peaks  is higher  for  the pristine samples  than for  the post-oxidized  ones,  indicating that  those signals are from the bulk material  beneath the scale layer. Additionally, the peak of MoSi2 phase  is absent in the both post-oxidized composites, a trace amount  of SiC was detected only in the post-oxidized ZMS20. This  indicates that SiC and MoSi2 are thoroughly oxidized during  the exposure. The crystalline phases identified in the as-sintered  and post-oxidized samples are summarized in Table 1.  It is known that ZrB2, MoSi2 and SiC phases oxidize to form  ZrO2, B2O3, SiO2 and MoB when ZrB2-SiC and ZrB2-MoSi2  composites are exposed to air at high temperature. Previous  studies in the ZrB2-SiC composition show that ZrB2 and SiC oxidize in air at 1500 8C, according to the following reactions  [8,10]: ZrB2 ðsÞ þ 2:5O2 ðgÞ ¼ ZrO2 ðsÞ þ B2O3 ðgÞ  (1)  SiCðsÞ þ 1:5O2 ðgÞ ¼ SiO2 ðsÞ þ COðgÞ  (2)  For the ZrB2-MoSi2 composition, on the other hand, ZrB2 and  MoSi2 phases oxidize  as  they were  exposed to air  at high  temperatures, according to the following reactions [12,17], MoSi2 ðsÞ þ 3:5O2 ðgÞ ¼ MoO3 ðgÞ þ 2SiO2 ðsÞ  (3)  ZrB2 ðsÞ þ 2MoSi2 ðsÞ þ 5O2 ðgÞ ¼ ZrO2 ðsÞ þ 2MoBðsÞ þ 4SiO2 ðsÞ  (4Þ  as a result of  the outward diffusion of constituent elements  cations from the bulk to the oxidized surface and the inward  diffusion of O through the scale layer. In the present study, XRD  analysis showed the presence of SiO2, ZrO2 and MoB phases in  the post-oxidized samples (Fig. 3). The absence of B2O3 phase  (f)  ZrB  MoSi  ZrO  SiO  MoB  SiC  ZMS20, 10h  ZrSiO  (d)  (e)  ZMS5, 10h  ZMS20, 1h  u  .  )  (b)  (c)  ZMS20  ZMS5, 1h  I  n  t  n e  s  i  t  y  (  a  .  (a) (a)  80  70  60  50  40  30  20  ZMS5  2 θ  (degree)  Fig. 3. X-ray diffraction patterns of the specimen surfaces for the ZrB2-MoSi2- SiC composites before and after oxidation exposure at 1500 8C for 1 h and 10 h:  (a and b) pristine and (c-f) post-oxidized.  60  2  /  c  m  4  )  40  50  1500 ° C  ZMS5  ZMS20  2  (  m  g  30  W  e  i  h g  t  G  a  i  n  ,  W  10  20  Time at Temperature, t (h)   0  12  10  8  6  4  2  0  Fig. 2. Parabolic plots of specific weight gain as a function of ZrB2-MoSi2-SiC composites oxidized at 1500 8C.  time for  the  Table 1  Crystalline phases in the pristine and oxidized ZrB2-MoSi2-SiC composites.  Materials  Before oxidation  After oxidation  1 h  10 h  ZMS5  Main phase  ZrB2, MoSi2  ZrO2, ZrB2  ZrO2  Minor phase  SiC  MoB, ZrSiO4  ZrB2, SiC, MoB, ZrSiO4  ZMS20  Main phase  ZrB2, MoSi2  ZrO2, ZrB2  ZrO2  Minor phase  SiC  ZrB2, MoB, ZrSiO4  ZrB2, SiC, MoB, ZrSiO4, SiO2  S. Guo et al. / Ceramics International 37 (2011) 585-591  587  [ ( ) T D $ F I G ]   [ ( ) T D $ F I G ]     \\x0c', 'is attributed to its vaporization above 1300 8C, as a result of its  high vapor pressure. Thus,  it could be expected that similar 1500 8C for  reactions  occurred  during  the  exposure  at  the  materials investigated in this study.  Thermodynamically,  although the  above-mentioned reac tions  are  favored  at  high  temperature,  they  did  not  occur  simultaneously and their occurrence depended on exposure 1500 8C,  temperature. At  the  reaction  (1)  has  the more  favorable followed by reaction (2), reactions (3) and (4) [18-  20]. Thus, when the ZrB2-MoSi2-SiC composite was exposed at 1500 8C,  to air  the oxidation  reaction conducted by the  following sequence:  (i) firstly ZrB2 oxidized by air oxygen  according to the reaction (1); (ii) then SiC oxidized according to  the  reaction  (2);  and  (iii)  finally MoSi2  oxidized  by  the  reactions (3) and (4). At the initial stage of exposure ZrB2 firstly  oxidized  into  ZrO2  and  B2O3,  resulting  in  deviation  of  oxidation from parabolic behavior and the rapid increase of  weight gain. Subsequently, SiC oxidized prior to MoSi2 and a  continuous and/or partially amorphous  silica-rich glass  scale  was produced on the surface of samples. The formation of the  scale,  in  particular  continuous  scale,  inhibited  the  inward  diffusion of oxygen in air, therefore improvement of oxidation  resistance. Consequently, the oxidation behavior was controlled  by inward diffusion of oxygen through the  scale  layer  and  outward diffusion of constituent elements cations  from bulk  materials. The later oxidation reaction of MoSi2 accelerated the  formation of the amorphous silica scale and increased the thick  of the scale. Thus,  it should be reasonable that  the addition of  SiC improved oxidation resistance of MoSi2-containing ZrB2  and this improvement was enhanced with increasing amount of  SiC added. After ZrB2, SiC and/or MoSi2 oxidized,  the ZrO2  and amorphous  silica coexisted in the silica-rich glass  scale  layer on the surface of samples. This results in a further reaction  between ZrO2 and amorphous  silica. An early study [21]  in  ZrO2  and  amorphous  silica  showed  that  interstitial  silicon  diffuses and dissolves into crystalline ZrO2 until  the solution  limit  is  reached when ZrO2 and amorphous  silica coexisted,  thereafter precipitation of ZrSiO4. In the present study, the peak  of ZrSiO4 was detected in the post-oxidized samples for both  compositions materials. For comparison, crystalline SiO2 was  detected only in the ZMS20 sample oxidized for 10 h. The  formation of crystalline SiO2 is not well-understood but related  with Si concentration in amorphous silica.  It seems to expect  the precipitation of  crystalline SiO2  from amorphous  silica  when Si concentration reached saturation.  Fig. 4 shows the typical surface morphologies for  the two  composites after oxidation. After oxidation of 1 h, the surfaces  are covered with a continuous silica layer where white ZrO2 particles of <1 mm diameter are embedded. For ZMS5,  the  ZrO2  particles  have  not  coarsened  significantly  even  after  oxidation of 10 h. For comparison, for ZMS20 after oxidation  of  10 h  the  particles  significantly  coarsened  to  form the  nodules. EDX analysis reveals that  the background consists of  only Si and O (Fig. 4(e)). This indicated that the background is  an amorphous silicate phase for both the composites. This is  similar to SiC-free ZrB2-MoSi2 composites oxidized under the  same oxidation conditions [17]. For comparison,  the nodules  particles were oxide cluster where some larger ZrO2 embedded  in a SiO2-rich glass matrix, verified by EDX analysis (Fig. 4(f)).  In  addition,  some  cracks  are  observed  around  the  nodules  (indicated by arrows in Fig. 4(d)). This cracking behavior  is  associated  either with  different  thermal  expansion  between  ZrO2 and SiO2 or with the volumetric change accompanying  transformation of cristobalite from bto a-phase as well as the  phase  transformation  of  tetragonal  to monoclinic ZrO2  on  cooling [22,23].  In Fig. 5,  the EPMA back-scattered images of  the cross section and corresponding X-ray image of various elemental  Fig. 4. Typical SEM images of the surface morphologies for the ZrB2-MoSi2-SiC composites oxidized at 1500 8C for 1 h (a and c) and 10 h (b and d), with EDX  spectra of  (e) background and (f) nodules particles.  S. Guo et al. / Ceramics International 37 (2011) 585-591  588  [ ( ) T D $ F I G ] \\x0c', 'mappings of  the composites oxidized for 10 h are presented.  The cross-section of the post-oxidized samples is divided into  the oxidized reactive region and the unreactive bulk material  region. For ZMS5, the thickness of the reactive region is equal to \\x18148 mm which is almost the same with that of SiC-free 40 vol.% MoSi2-containing ZrB2 oxidized at 1500 8C for 10 h  [17]. This  reveals  that  5 vol.% SiC added  is  significantly  ineffective for  further  improving oxidation resistance of  the  ZrB2-MoSi2  composition.  For  ZMS20, the thickness of to \\x1862 mm which is much  oxidized reactive region is equal  thinner  than that of ZMS5. Thus, 20 vol.% SiC added further  significantly improved the oxidation resistance. In addition, the  oxidized reactive region consists of two different characteristic  layers, I and II. Layer I is a dense layer and it is rich Si and O  (Fig. 5(b) and (d)). The thickness of layer I is much thicker than  for ZMS20 than for ZMS5. This means that 20 vol.% SiC added  formed a thicker dense scale layer which prevented effectively  the  inward  diffusion  of  oxygen  through  it,  therefore  high  oxidation resistance. Although the thickness of  layer  I  in the  studied materials is much thinner than for the SiC-free 40 vol.%  MoSi2-containing ZrB2,  the oxidation resistance is higher  for  the former than for the latter [17]. Obviously, this suggests that  the oxidation resistance is not only linked with thickness of the  scale layer, but also with its viscosity and compositions.  Moreover,  layer  II  strongly  depends  on  amount  of SiC  added, being much thicker  for ZMS5 than for ZMS20. This  suggests  the  layer  I  is more  effective  for  inhibiting inward  diffusion  of  oxygen  during  exposure  for ZMS20  than  for  ZMS5. This is associated with the presence of some crystalline  SiO2 phase in it because layer I contains some crystalline SiO2  phase for ZMS20 but the absence of crystalline SiO2 phase for  ZMS5, with a thicker dense layer I for the former as well. Layer  II  is very complex and it contains O, Si, Mo, Zr and B. The  morphology of layer II is strongly dependent on the amount of  SiC added. For ZMS5,  it  is found layer upon layer of layer II  which is  stacked by (Si, O)-rich layer and (Mo, Zr, B)-rich  layer. A similar structure was previously reported in 40 vol.% MoSi2-containing ZrB2 oxidized at 1500 8C for 10 h [17]. For  comparison, in the case of ZMS20, layer II consists of a single  layer. The unreactive bulk material region contains Si, Mo, Zr,  C and B, but no O. Hence, the unreactive bulk material consists  of Zr-B, Mo-Si and Si-C phases.  3.3. Compositions and distribution of reactive compounds  Fig. 6 shows the detailed oxidation products distribution and  compositions in the cross-section for the composites oxidized  for 10 h, with the scatter diagram method under EPMA. The Si-  O phase is the major reactive compound in layer I for the two  compositions, and a trace amount of Zr-O and M-B phases is  present in it as well. Layer II has the complex compounds and it  consists of Si-O, Mo-B, Zr-O, and Si-C, Zr-B phases. For  ZMS5,  layer  II  is divided into two sublayers:  II(a) and II(b).  Sublayer  II(a) consists of Si-O, Mo-B, and Zr-O phases  in  which Zr-O is the primary phase, while SiC, ZrB2 and MoSi2  were not shown. In sublayer II(b),  the Zr-B phase and a trace  amount of Si-C phase are also present, but the Mo-Si phase is  absent  in the both layers. Differing from ZMS5,  in ZMS20  sublayer  II(a)  is  absent, only sublayer  II(b)  is observed.  In  sublayer II(b), the Zr-O, Mo-B and SiC phases are present, and  Fig. 5. Typical examples of EPMA backscattered image of the cross-section and elemental mappings under EPMA for (a and b) ZMS5 and (c and d) ZMS20 oxidized at 1500 8C for 10 h.  S. Guo et al. / Ceramics International 37 (2011) 585-591  589  [ ( ) T D $ F I G ] \\x0c', '590  S. Guo et al. / Ceramics International 37 (2011) 585-591  Fig. 6. Examples of X-ray images of phase mapping of the cross-section under EPMA and the scatter diagrams of Si vs. O, Zr vs. O, Mo vs. B, Mo vs. Si, Si vs. C and (a) ZMS5 and (b) ZMS20 oxidized at 1500 8C for 10 h.  Zr vs. B for  a trace amount of Si-O phase is present as well. Thus, layer II is  diagrams of the obtained characteristic X-ray intensities for Si-  a SiCand MoSi2-depleted zone where the oxidation reaction  O, Zr-O, Mo-B, Mo-Si, Si-C, and Zr-B phases are also shown  occurred,  and it  is more  extended for ZMS5,  compared to  in Fig. 6, in which each point represents the number of the same  ZMS20. Hence,  the  addition of SiC promoted formation of  intensities. Note that  the solid circle in each figure represents  thicker dense glassy scale on the surface of  samples which  the pure  stoichiometric SiO2, ZrO2, MoSi2, SiC,  and ZrB2  inhibited effectively outward diffusion of oxygen through it,  phases. It is found that the crystalline Zr-O, M-B, Mo-Si, Si-C  therefore improving oxidation resistance. The unreactive bulk  material  region is  the same for  the two compositions and it  consists of Zr-B, Mo-Si and Si-C phases.  and Zr-B phases are the stoichiometric ZrO2, MoB, MoSi2, SiC  and ZrB2 phases for both compositions. On the other hand, for  ZMS5 the composition of amorphous Si-O phase is determined  In order to identify the reactive compound compositions as  to be 33.2 mass.% Si and 52.0 mass.% O, deviating from the  well as to examine a change of the chemical compounds during  pure SiO2 phase of 46.74 mass.% Si and 53.26 mass.% O. This  exposure, the analysis of X-ray image data for Si, O, Zr, Mo, C  means that Si  is depleted in outermost amorphous scale, as a  and B elements is conducted on the cross-section. The scatter  result of diffusion and dissolution of Si into ZrO2 [21]. Similar  [ ( ) T D $ F I G ] \\x0c', 'behavior was previously reported in SiC-free 40 vol.% MoSi2 containing ZrB2 composites [17]. Differing with the previous  study,  the nonstoichiometric MoSi2 was not detected in this  study.  This  shows  that  although  5 vol.% SiC  added  is  ineffective for significantly improving oxidation resistance of  40 vol.% MoSi2-containing ZrB2, but  inhibiting oxidation of  MoSi2  in the unreactive bulk materials. For comparison,  for  ZMS20 the Si-O phase is determined to be 41.9 mass.% Si and  53.8 mass.% O, being close with a pure SiO2 phase. This  suggests that 20 vol.% SiC added supplies amount of Si needed  for forming a pure SiO2 phase. Based on the above-mentioned  amorphous silica phase composition,  it seems to be expected  that the crystalline SiO2 phase in ZMS20 precipitated from the  amorphous silica phase which composition was close to pure  SiO2 on cooling, as a result of saturation of Si concentration. In  addition, analysis of the scatter diagrams of Zr-B, Mo-Si and  Si-C shows that these compounds are stoichiometric. Previous  study of ZrB2-MoSi2 shows that Mo-Si phase is nonstoichio metric MoSi2, as a result of the outward diffusion of Si during  exposure. Thus,  the addition of SiC seems to be effective for  supplying enough Si source for precipitating a crystalline SiO2  phase from amorphous silica as well as increasing viscosity of  the amorphous silica-rich glass scale layer, therefore inhibiting  the outward diffusion of Si  from the unreactive bulk and the  inward  diffusion  of  oxygen  through  the  scale  layer  during  oxidation exposure.  4. Conclusions  In conclusion, the oxidation resistance of the ZrB2-MoSi2-  SiC composites  is  improved with  SiC addition,  and  the  improvement  enhanced with  amount  of  SiC added.  The  microstructure of  the oxidized ZrB2-MoSi2-SiC composites  consists of an outermost dense glassy scale, middle oxidized  reactive layer, and an unreactive bulk material. The outermost  dense scale layer is much thicker for ZMS20 than for ZMS5.  However,  the middle oxidized reactive layer is much greater  for ZMS5 than for ZMS20. The dense scale layer consists of  ZrO2 and amorphous silica where ZrO2 embedded in SiO2 rich glass matrix, a trace amount of MoB and ZrSiO4 was  presented. For ZMS20,  a  crystalline SiO2  phase was  also  present  in the post-oxidized sample  for 10 h, with a  trace  amount of SiC. The middle  reactive  layer  is  composed of  complex compounds, and it consisted of amorphous SiO2-rich  glass,  crystalline  ZrO2, MoB,  SiC  and  ZrB2  for  both  compositions. The unreactive bulk material  is  the same for  both compositions  composites,  consisting of ZrB2, MoSi2,  and SiC.  Acknowledgements  The authors would like to thank Dr. T. Kimura and Mr. T.  Aoyagi, National  Institute  for Materials  Science,  for  his  assistance with EPMA measurements as well as discussion.  References  [1] C. Mroz, Zirconium diboride, Am. Ceram. Soc. Bull. 73 (1994) 141-  142.  [2] K. Upadhya, J.M. Yang, W.P. Hoffmann, Materials for ultrahigh temper ature structural applications, Am. Ceram. Soc. 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Fahrenholtz, G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500 8C, J.  Eur. Ceram. Soc. 27 (2007) 2495-2501.  [11] D. Sciti, M. Brach, A. Bellosi, Long-term oxidation behavior and me chanical  strength degradation of a pressurelessly sintered ZrB2-MoSi2  ceramic, Scripta Mater. 53 (2005) 1297-1302.  [12] D. Sciti, M. Brach, A. Bellosi, Oxidation behavior of  a pressureless  sintered  ZrB2-MoSi2  ceramic  composite,  J. Mater. Res.  20  (2005)  922-930.  [13] S.Q. Guo, T. Nishimura, T. Mizuguchi, Y. Kagawa, Mechanical properties  of  hot-pressed ZrB2-MoSi2-SiC composites,  J. Eur. Ceram. Soc.  28  (2008) 1891-1898.  [14] S.Q. Guo, T. Nishimura, H. Tanaka, Y. Kagawa, Thermal and electrical  properties  in hot-pressed ZrB2-MoSi2-SiC composites,  J. Am. Ceram.  Soc. 90 (2007) 2255-2258.  [15] T. Kimura, K. Nishida, S. Tanuma, Spatial  resolution of a wavelength dispersive electron probe microanalyzer equipped with a thermal field  emission gun, Microchim. Acta 155 (2006) 175-178.  [16] T. Kimura, T. Sugizaki, K. Nishida, N. Ishikawa, S. Tanuma, Analysis of  joining boundary between Ni-P Electroless plate and solder by EPMA  scatter diagram method, J. Jpn Inst. Metals. 68 (2004) 8-13.  [17] S.Q. Guo, T. Mizuguchi, T. Aoyagi, T. Kimura, Y. Kagawa, Quantitative  electron microprobe characterizations of oxidized ZrB2 containing MoSi2  additives, Oxid. Metals 72 (2009) 335-345.  [18] W.G. Fahrenholtz, G.E. Hilmas,  I.G. Talmy,  J.A. Zaykoski, Refractory  diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-  1364.  [19] Z.H. Yang, D.C.  Jia, Y. Zhou, P.Y. Shi, C.B. Song, L. Lin, Oxidation  resistance of hot-pressed SiC-BN composites, Ceram. Int. 34 (2) (2008)  317-321.  [20] Y.Q. Liu, G. Shao, P. Tsakiropoulos, On the oxidation behavior of MoSi2,  Intermetallics 9 (2) (2001) 125-136.  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  "_id": 152,
  "PDF": "Oxidation Behavior of ZrB2–YAG–Al2O3Ceramics at High Temperature.pdf",
  "Text": "['Materials and Manufacturing Processes  ISSN: 1042-6914 (Print) 1532-2475 (Online) Journal homepage: https://www.tandfonline.com/loi/lmmp20  Oxidation Behavior of ZrB2-YAG-Al2O3 Ceramics at High Temperature  Jie-Guang Song  To cite this article: Jie-Guang Song (2010) Oxidation Behavior of ZrB2-YAG-Al2O3 Ceramics at High Temperature, Materials and Manufacturing Processes, 25:8, 724-729, DOI: 10.1080/10426910903229362  To link to this article:  https://doi.org/10.1080/10426910903229362  Published online: 02 Sep 2010.  Submit your article to this journal   Article views: 70  View related articles   Citing articles: 4 View citing articles   Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=lmmp20  \\x0c', 'Materials and Manufacturing Processes, 25: 724-729, 2010  Copyright © Taylor & Francis Group, LLC ISSN: 1042-6914 print/1532-2475 online DOI: 10.1080/10426910903229362  Oxidation Behavior of ZrB2-YAG-Al2O3 Ceramics at High Temperature  Jie-Guang Song1(cid:1)2  1 School of Mechanical and Materials Engineering, Jiujiang University, Jiujiang, China 2 Jiujiang Key Laboratory of Green Remanufacturing, Jiujiang University, Jiujiang, China  ZrB2 , YAG, and Al2O3 are widely applied because of some excellent performance qualities, but ZrB2 is easily oxidized in high-temperature air. To make the ZrB2 ceramics obtain better oxidation resistance, high-density ZrB2 -YAG-Al2O3 ceramics were prepared. The oxidation behavior of ZrB2 -YAG-Al2O3 ceramics was investigated. The weight gain of ceramics was increased with increasing oxidation temperature, and the weight gain of ceramics was reduced with increasing YAG-Al2O3 content and Al2O3 proportion, especially above 1500 increased with prolonged oxidation time, but the rate of weight gain was decreased gradually after 2 h, and the weight gain was reduced with increasing YAG-Al2O3 content and Al2O3 proportion. Increasing YAG-Al2O3 content and Al2O3 proportion may raise the oxidation resistance of ZrB2 -YAG-Al2O3 ceramics. Through the above analysis, the oxidation resistance model of ZrB2 -YAG-Al2O3 ceramics was constructed to fully comprehend the oxidation process and mechanism.  C. The weight gain was  (cid:1)  Keywords High temperature; Oxidation; Weight gain; ZrB2 ceramics.  Introduction  Zirconium diboride (ZrB2 ) has attracted substantial interest because of its extreme chemical and physical properties, such as high melting point, superior hardness, and low electrical resistance. ZrB2 has several applications such as Hall-Heroult cell cathodes for electrochemical processing of aluminum, evaporation boats, crucibles for handling molten metals, thermowell tubes for steel refining, thermocouple sleeves for high-temperature uses, nozzles, plasma electrodes, or as dispersoids in metal and ceramic-matrix composites for heaters and igniters [1-4]. However, ZrB2 is easily oxidized in high-temperature air, which impacts high-temperature strength and restricts its application range [5, 6]. Some excellent oxidation resistance materials are considered as the assistant phase of ZrB2 materials to prepare ZrB2 composite materials for improving the high-temperature performance of ZrB2 materials, such as Al2O3-ZrB2 , SiC-ZrB2 , LaB2 -ZrB2 , and ZrO2 -ZrB2 [7-9]. Yttrium aluminum garnet (YAG or Y3Al5O12 ) adopts a cubic garnet structure and is of great interest as a high-temperature engineering materials due to its hightemperature strength coupled with low creep rates [10-14], which indicates that YAG ought to be a suitable matrix or reinforcing material [15, 16]. To make ZrB2 ceramics obtain better oxidation resistance, high-density ZrB2 -YAG-Al2O3 ceramics were prepared. In this article, the oxidation behavior of ZrB2 -YAG-Al2O3 ceramics is investigated. Because the weight of ZrB2 -YAG- Al2O3 ceramics is varied during the oxidation process, the  oxidation degree of ZrB2 -YAG-Al2O3 ceramics may be characterized through weight gain of ZrB2 -YAG-Al2O3 ceramics below a certain temperature. The weight gain of ZrB2 -YAG-Al2O3 ceramics is discussed to investigate the oxidation behavior of ZrB2 -YAG-Al2O3 ceramics.  Materials and experiment  (cid:1)  D50  Analytical-grade aluminum nitrate, yttrium nitrate, ammonia, and commercially available ZrB2 powder (99.5% purity, 12 (cid:2)m) were used. ZrB2 particles were coated with Al2O3-Y2O3 composite particles via the coprecipitation method [17]. Different types of ZrB2 -YAG- Al2O3 ceramics (relative density of 95%) were prepared using Al2O3-Y2O3 /ZrB2 shell-core composite particles with spark plasma sintering (SPS) under sintering temperature of 1700 C, holding time of 4 min, and sintering pressure of 20 MPa (Table 1). ZrB2 -YAG-Al2O3 ceramics were oxidized by air in the furnace, and the weight and microstructure of ZrB2 -YAG-Al2O3 ceramics were tested after oxidation. The process flow diagram is shown in Fig. 1. ZrB2 -YAG-Al2O3 ceramics were prepared using SPS (model SPS-1050, Japan). Oxidation treated in the furnace (model Nabertherm LHT04, Germany). Phase analysis was identified by X-ray powder diffraction (XRD; model D/Max-RB, Japan). Microstructure analysis was performed by scanning electron microscopy (SEM; model: JSM5610LV, Japan). Element analysis was performed with an electron probe microscopy apparatus (EPMA; model: JXA880R, Japan).  Received June 14, 2009; Accepted August 2, 2009 Address correspondence to Jie-Guang Song, School of Mechanical and Materials Engineering, Jiujiang University, Jiujiang 332005, China; E-mail: songjieguang@yahoo.com.cn  Effect of Oxidation Temperature on Oxidation Resistance of ZrB2 -YAG-Al2O3 Ceramics  The weight gain of different ceramics after oxidized under different oxidation temperature for 1 h is shown in Fig. 2,  Results and discussion  724  \\x0c', 'OXIDATION OF CERAMICS AT HIGH TEMPERATURE  725  Table  1.—Different  types  of  ceramics  examined  in  the  current work.  Ceramic  Z-YA Z-Y3A Z-Y6A  YAG:A12 O3  (mol)  1:1 1:3 1:6  Phase  ZrB2 + YAG + A12O3 ZrB2 + YAG + A12O3 ZrB2 + YAG + A12O3  which indicates that the weight gain of all sorts of ceramics is increased with increasing oxidation temperature, and the weight gain of ceramics is reduced with increasing YAG- Al2O3 content and Al2O3 proportion, especially above 1500 C. The effect of oxidation temperature on the weight gain of ceramics is decided by the relative density of ceramics. However, factors effect of YAG-Al2O3 content and YAG: Al2O3 on the weight gain of ceramics are not relative density of ceramics but also the chemical reaction, which is ZrB2 reacted with O2 to form B2O3 [18], and B2O3 is reacted with Al2O3 to form Al18B4O33 (Fig. 3). Al18B4O33 is melted and coated on the surface of ceramics to form a protective layer for the oxidation resistance of ceramics at high temperature [19] (Fig. 4). XRD of ceramics with different Al2O3 content after oxidation at 1600 C for 1 h is shown in Fig. 5, which indicates that the Al18B4O33 content of the oxidation surface of ceramics is increased with increasing Al2O3 proportion, which makes the thickness of the oxidation resistance layer increase. The effect of oxidation temperature on the oxidation layer thickness of Z-Y6A ceramics is shown in Fig. 6, which indicates that the oxidation layer thickness is increased with increasing oxidation temperature, and the oxidation layer thickness is reduced with increasing Y6A content, especially 1500 C. The microstructure of the oxidation layer thickness  (cid:1)  (cid:1)  (cid:1)  Figure 2.—Effect of oxidation temperature on weight gain of ceramics after oxidation of 1 h: (a) Z-YA, (b) Z-Y3A, and (c) Z-Y6A.  after oxidation under different oxidation temperatures is shown in Fig. 7, which indicates that the oxidation layer is loosen.  Effect of Oxidation Time on Oxidation Resistance of ZrB2 -YAG-Al2O3 Ceramics  The weight gain of ZrB2 -YAG-Al2O3 ceramics at 1200 C for different oxidation times is shown in Fig. 8,  (cid:1)  Figure 1.—Process flow diagram.  \\x0c', '726  J.-G. SONG  Figure 3.—Effect of oxidation temperature on phase of Z-40wt%YA ceramics after oxidation of 1 h: (a) 1100 C, (b) 1200 C, (c) 1300 C, (d) 1400 C, (e)  (cid:1)  (cid:1)  (cid:1)  (cid:1)  (cid:1)  1500  C, and (f) 1600  C.  (cid:1)  Figure 6.—Effect of oxidation temperature on thickness of oxidation layer of Z-Y6A ceramics after oxidation for 1 h.  Figure 4.—Element distribution of ZrB2 -40wt%Y6A ceramics after oxidation at 1600 C for 1 h.  (cid:1)  Figure 5.—XRD of surface of different ceramics after oxidation at 1600 for 1 h: (a) Z-40wt%YA, (b) Z-40wt%Y3A, and (c) Z-40wt%Y6A.  (cid:1)  C  Figure  7.—Effect  of  oxidation  temperature  on  oxidation  40wt%Y6A ceramics  after  oxidation  of  1 h:  (c) 1300  C, (d) 1400  C, (e) 1500  C, and (f) 1600  C.  (a)  (cid:1)  (cid:1)  (cid:1)  (cid:1)  (cid:1)  layer  C,  (b)  of ZrB2 1200 C,  (cid:1)  1100  \\x0c', 'OXIDATION OF CERAMICS AT HIGH TEMPERATURE  727  Figure 9.—Effect of oxidation time on thickness of oxidation layer of Z-Y6A ceramics at oxidation temperature of 1200 C.  (cid:1)  Figure 8.—Effect of oxidation time on weight gain of ceramics at oxidation temperature of 1200 C: (a) Z-YA, (b) Z-Y3A, and (c) Z-Y6A.  (cid:1)  which indicates that the weight gain is increased with prolonged oxidation time, but the rate of weight gain is decreased gradually after 2 h, and the weight gain is reduced  Figure 10.—Effect of oxidation time on oxidation layer of Z-40wt%Y6A ceramics at oxidation temperature of 1200 C: (a) 1 h, (b) 2 h, (c) 3 h, (d) 4 h,  (cid:1)  (e) 5 h, and (f) 6 h.  \\x0c', '728  with increasing YAG-Al2O3 content and Al2O3 proportion. Because more B2O3 is produced with prolonged oxidation time, B2O3 is melted and coated on the surface of ceramics to protect the ceramics. The relative density of ceramics is increased with increasing YAG-Al2O3 content and Al2O3 proportion, which exists few channel for O2 diffusion in the ceramics [20]. The oxidation layer thickness of Z-Y6A ceramics after oxidation at 1200 C for different oxidation times is shown in Fig. 9, which indicates that the oxidation layer thickness is increased with prolonging oxidation time (Fig. 10), and the oxidation layer thickness is reduced with increasing YAG-Al2O3 content and Al2O3 proportion.  (cid:1)  Model of Oxidation Resistance of ZrB2 -YAG-Al2O3 Ceramics  (cid:1)  Through the above analysis and results, the oxidation model of ZrB2 -YAG-Al2O3 ceramics at high temperature is constructed, which helps to understand the oxidation mechanism of ZrB2 -YAG-Al2O3 ceramics. A sketch of the oxidation model of ZrB2 -YAG-Al2O3 ceramics is shown in Fig. 11. Under 1200 C, the relative density of the oxidation layer (Fig. 11(a), district) is higher due to the existing YAG and Al2O3 phase, which remains the few channel for O2 diffusion to react with ZrB2 (Fig. 11(a), district) in the ceramics, and the ceramics (Fig. 11(a), district) are protected to reach the aim of oxidation resistance. Above 1200 C, ZrB2 is reacted with O2 to form B2O3 (Fig. 11(b), district), and B2O3 is reacted with Al2O3 to form Al18B4O33 (Fig. 3). Al18B4O33 is melted and coated on the surface of ceramics to form a protective layer for the oxidation resistance of ceramics at high temperature (Fig. 11(b), district).  (cid:1)  Figure 11.—Oxidation model of ceramics: 1200 C. : Ceramics; : ZrO2 ; : YAG;  (cid:1)  (a) below 1200 : Al2O3 .  (cid:1)  C and (b) above  J.-G. SONG  Conclusion  (cid:1)  The weight gain of ceramics is increased with increasing oxidation temperature, especially above 1500 C. The weight gain of ceramics is reduced with increasing the YAG-Al2O3 content and Al2O3 proportion. The weight gain is increased with prolonged oxidation time, but the rate of weight gain is decreased gradually after 2 h, and the weight gain is reduced with increasing YAG-Al2O3 content and Al2O3 proportion. Increasing YAG-Al2O3 content and Al2O3 proportion may raise the oxidation resistance of ZrB2 - YAG-Al2O3 ceramics. Through the above analysis, an oxidation resistance model of ZrB2 -YAG-Al2O3 ceramics was constructed to fully comprehend the oxidation process and mechanism.  Acknowledgments  The authors are thankful for the financial support provided by the Science Fund for Young Scholars of the Educational Department of Jiangxi Province, China (Grant No. GJJ09595) and for the apparatus provided by the Center for Materials Testing and the Functionally Gradient Material (FGM) Research Laboratory of Wuhan University of Technology, China.  References  459,  1. 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},{
  "_id": 153,
  "PDF": "Oxidation behavior of ZrC-based composites in static laboratory air up to 1300°C.pdf",
  "Text": "[\"Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  Contents lists available at ScienceDirect  Int. Journal of Refractory Metals and Hard Materials  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / I J R M H M  Oxidation behavior of ZrC-based composites in static laboratory air up to 1300 °C  Baoxia Ma a,⁎, Wenbo Han b, Erjun Guo a  a Department of Materials Science and Engineering, Harbin University of Science and Technology, Harbin 150040, China b Center for Composite Materials, Harbin Institute of Technology, Harbin 150001,China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 14 January 2014 Accepted 27 May 2014 Available online 2 June 2014  Keywords:  ZrC SiC Ceramic composites Oxidation  Introduction  Two ZrC-based composites, viz., ZrC + 20 vol.%SiC and ZrC + 20 vol.%SiC + 20 vol.%ZrB2, were prepared by hot pressing. The oxidation behavior of the composites was studied in the temperature range of 800 °C-1300 °C in air. From the mass gain tests, the surface and cross-sectional SEM observation and EDS analysis of the oxidized samples, it was found that the ZrC + 20 vol.%SiC composite seemed to have a high oxidation resistance below 1000 °C, but the ZrC + 20 vol .%SiC + 20 vol.%ZrB2 composite appeared a more excellent oxidation resistance up to 1200 °C. It is main reason that the oxidation rate of ZrC is higher than that of SiC below 1300 °C, meaning that SiO2 from SiC is insufficient to form a dense tight protection film; however, the addition of ZrB2 particles promotes the formation of more borosilicate above 1000 °C, which exerts more protective effect by self-healing mechanism.  © 2014 Elsevier Ltd. All rights reserved.  U ltra h igh-temperature ceram ics (UHTCs) , such as carbides and borides of Zr and Hf , are widely investigated as candidates for high-temperature structural applications that require exposure to environments over 3000 °C because of their unrivaled melting points exceed ing 3000 °C , coup led to retained strengths and a therma l stability at very high temperatures [1-3]. Recently, ZrC as carbidesbased ceramics receives extensive attention due to its superiority in higher melting point, lower manufacturing costs and more excellent thermal shock resistance and high-temperature strength compared with borides-based ceramics. Various sintering methods such as hot press ing [4] , pressureless s intering [5] and hot isostat ic pressing [6] are employed to achieve ZrC ceramic composites with high densification density and high strength. Moreover , the introduction of some additives can decrease significantly the sintering temperature and improve greatly the fracture toughness, in addition to obtain better densification. Considered that in high-temperature applications ZrC structural components are inevitably exposed to elevated temperatures in oxidizing environments, the oxidation of ZrC under these conditions plays a key role in influencing the material's performance. Therefore, the oxidation behavior of ZrC ceramic materials needs to be investigated. Research on oxidation of monophase ZrC ceramic has been reported  ⁎  Corresponding author. Tel.: + 86 451 8639 2517. E-mail addresses: qiufeng-@126.com, mabaoxia@126.com (B. Ma).  http://dx.doi.org/10.1016/j.ijrmhm.2014.05.013 0263-4368/© 2014 Elsevier Ltd. All rights reserved.  ever since 1960s [7-9]. These several works on oxidation of ZrC have focused on the reaction at temperatures below 800 °C, and the description of oxidation mechanism mainly from kinetic viewpoints. In subsequent studies by Shimada et al. [10,11], the oxidation at various oxygen pressures and relatively high temperature (b 1000 °C) was performed, and the form of carbon atoms in the oxide product phases was reported. Although great efforts have been made in studying the oxidation of ZrC ceramics, so far few works has been carried out on the oxidation resistance of ZrC-based composites above 800 °C. The present article aimed to evaluate the oxidation behavior of two different materials (ZrC-SiC and ZrC-SiC-ZrB2 composites) at temperature up to 1300 °C in air. The microstructural evolution was analyzed on the surface, and the cross section of the oxidized samples and oxidation mechanisms was discussed.  Experimental  Commercial raw materials of ZrC (N 98%, mean particle size (FSSS) of 1.25 μm, Changsha Wing High High-Tech New Material Co., Ltd, China), SiC (99.5%, mean particle size of 2 μm, Weifang Kaihua Micro-powder Co., Ltd., China) and ZrB2 (N 99.4%, a mean size of 2 μm, Northwest Institute for Non-ferrous Metal research, China) were used to produce the ceramic composites. Two powder mixtures were prepared, i.e., ZrC + 20 vol.%SiC and ZrC + 20 vol.%SiC + 20 vol.%ZrB2, and corresponding composites were labeled as ZrC + SiC and ZrC + SiC + ZrB2, respectively. The powder mixtures were wet-milled for 12 h in a polyethylene bottle using ZrO2 balls and ethanol as milling media, subsequently dried in a rotary evaporator to minimize segregation. The obtained powder  \\x0c\", '160  B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  mixtures were placed in a graphite die and then hot pressed at 1900 °C for 60 min under a uniaxial load of 30 MPa in Ar atmosphere. Billets with a diameter of 50 mm and thickness of 10 mm were produced. Microstructure and mechanical properties of the as-sintered ZrC + SiC and ZrC + SiC + ZrB2 materials were described in previous study [12]. The oxidation tests were carried out in a muffle furnace in laboratory air between 800 °C and 1300 °C for different time from 10 min to 90 min using rectangular blocks with dimensions 5.0 mm × 5.0 mm × 10.0 mm. The rectangular block samples were cut from sintered materials by electrical discharge machining method. Before oxidation, the samples were cleaned in ultrasonicated acetone bath, dried and weighed using electronic balances with accuracy 0.01 mg, after all the surfaces were mechanically ground and polished with diamond abrasives down to 1 μm finish. After oxidation, the samples were weighted again in order to measure the weight change. Oxidized surfaces and polished crosssections were analyzed by a scanning electron microscopy (SEM) (Model FEI Sirion-200, Holland) equipped with energy-dispersive XRay spectrometer (EDS) (Model Link-Isis, America). Cross sections for microstructural analysis were perpendicular to the top surface of the oxidized blocks and then polished to a 1 μm finish using diamond abrasives. The phase composition after oxidation was analyzed by Xray diffraction (XRD) (Rigaku, Dmax-rb, Japan).  Results and discussion  The mass variation during oxidation tests  Fig. 1 shows the mass variation with oxidation time for the ZrC + SiC and ZrC + SiC + ZrB2 composites at 1000 °C, 1200 °C and 1300 °C. It is clear from Fig. 1(a) that the mass gain per surface area of ZrC + SiC + ZrB2 composite exhibits a linear increase with increasing oxidation time during oxidation at 1000 °C and is higher than that of ZrC + SiC composite at the same time; conversely, the mass gain of ZrC + SiC composite is nonlinear and the increasing rate of mass change is low throughout the oxidation process, especially when the time exceeds 30 min, the mass change is less. This suggests that the oxidation resistance of ZrC + SiC composite is higher than that of ZrC + SiC  + ZrB2 composite at 1000 °C. At 1200 °C, during the first 10 min, a rapid mass gain of ZrC + SiC + ZrB2 composite occurs, and this mass gain is higher than that of ZrC + SiC composite. However, after oxidation for 30 min, the mass gain of ZrC + SiC + ZrB2 composite is followed by a reduction and is lower than ZrC + SiC composite. As the oxidation temperature increased to 1300 °C, the oxidation trend is opposite to 1000 °C, as shown in Fig. 1(c). The mass gain of ZrC + SiC + ZrB2 composite evidently decreases. After oxidation for 90 min, the mass gain of ZrC + SiC + ZrB2 composite is only 14.18 mg/cm2 and is one third of ZrC + SiC composite. Less main gain of ZrC + SiC + ZrB2 composite above 1200 °C is closely related to oxidation product. In addition, it is found that the variation law in the mass gain of ZrC + SiC + ZrB2 composite at 1300 °C is in accordance with the typical oxidation kinetic [13]. The relation of oxidation mass gain vs. time of ZrC + SiC + ZrB2 composite at this temperature can be described by the following equation [14]:  Δm ¼ kp \\x02 t x  ð1Þ  where Δm is the mass gain per unit area after oxidation, t is the oxidation time and kp and x are oxidation rate constants. According to Eq. (1), the fitting curve of main gain vs. time is shown in Fig. 1(d). It may be concluded that the oxidation behavior of ZrC + SiC + ZrB2 composite obeys the parabolic law. And the oxidation resistance of ZrC + SiC + ZrB2 composite is much more excellent at high temperature of 1300 °C, which can be further verified by the morphology of the surface and cross section after oxidation.  Macroscopical morphology during oxidation  Fig. 2 shows the macrographs of ZrC + SiC and ZrC + SiC + ZrB2 composites at different temperatures for different oxidation times. It can be seen that all of the oxidized surfaces of ZrC + SiC composite appear white, and white oxides thicken gradually with the increasing of oxidation times, which is accompanied by the occurrence of volume expansion. As the temperature increases, the tendency of expansion gradually increases, and the bulk oxides on the samples fall off after  Fig. 1. Mass gain per surface area of ZrC + SiC and ZrC + SiC + ZrB2 composites as a function of oxidation time in air at (a) 1000 °C, (b)1200 °C , (c) 1300 °C and (d) fitting of (c).  \\x0c', 'B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  161  Fig. 2. Images of ZrC composites after oxidation at different temperatures and times (a) ZrC + SiC and (b) ZrC + SiC + ZrB2.  oxidation above 1000 °C and black unoxidized ZrC + SiC composite matrix is exposed. It should be noted that the shedding of oxide layer occurs during cooling process of the samples that are removed from high-temperature furnace. Compared with ZrC + SiC composite (Fig. 2(a)) , it is found that ZrC + SiC + ZrB2 composite shows opposite oxidation result with increasing oxidation temperature and prolonging oxidation time. After oxidation at temperatures below 1000 °C , the shedding of surface oxides appears. However, when the oxidation temperature increases to above 1200 °C, under the same time conditions, the oxidation undergoes significant changes , namely , no surface oxides fall off and the oxide layer become thinner. The macroscopical volume expansion is not almost observed after oxidation at 1300 °C.  Microstructural evolution during oxidation  Surface morphology changes occurring to ZrC + SiC composite as a result of oxidation at different temperatures for 30 min are shown in Fig. 3. It can be clearly seen that ZrC + SiC composite displays light gray surface with dark gray dispersed particles, in which light gray areas are ZrO2 and dark gray particles are SiC or SiO2. In the temperature range 800 °C-1000 °C, because the oxidation rate of SiC is much slower than that of ZrC [15-17], the SiC particles do not oxidize, which is verified by EDS analysis in Fig. 3(b). However, zirconium and oxygen are detected by EDS. Possible reason is that the freshly formed ZrO2 in the oxide layer is carried and transmitted to the surface during the upward escaping of gaseous oxides, followed by depositing on the surface after  Fig. 3. SEM micrographs of the surfaces for ZrC + SiC composite after oxidation for 30 min at (a) 800 °C, (b) 1000 °C, (c) 1200 °C and (d) 1300 °C.  \\x0c', \"162  B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  cooling. As ZrC oxidizes, SiC particles are embedded in the growing ZrO2 layer. In addition, it is worth noting that some microcracks are visible on the sample surface, and they appear mainly at the interface between ZrO2 and dark gray particles. It is assumed that the generation of surface microcracks result from the thermal stresses due to mismatch of thermal expansion coefficients (αZrC = 6.7 × 10−6 oC−1 [18], αZrO2 = 10.5 × 10− 6 oC− 1 [19,20] and αSiC = 3.9 × 10− 6 oC− 1 [21]) during cooling process or growing stresses of ZrO2. In the samples oxidized at 1200 °C-1300 °C, as shown in Fig. 3(c) and (d), the surface morphology shows no significant changes except that the amount of surface microcracks reduces and the oxidation surface becomes flatter. It can be found from high magnification surface for ZrC + SiC composite in Fig. 4 that SiO2 glass phase emerges around initial SiC particles (as shown by arrows). According to Kobayashi et al. [22], SiC starts to be oxidized to SiO2 at about 900 °C, and the oxidation rate of SiC is much slow at temperatures below 1200 °C. Therefore, the oxidation of SiC is limited, and the amount of SiO2 glass is relatively less on the surface of samples oxidized at the temperature of b 1200 °C. When the temperature approaches and exceeds 1200 °C, SiO2 glass is readily observed on the surface of the samples. In this case, SiO2 glass can relax the stress and seal some microcracks to large extent. Fig. 5 shows SEM micrographs and EDS analysis of the surface of ZrC + SiC + ZrB2 composite after oxidation at different temperatures for 30 min. After oxidation at 800 °C, a few flaky and transparent substances that grow in laminated form are observed on the oxidation surface, which appears to be composed of a large number of oxygen and an extremely small amount of boron and zirconium by EDS analysis. This result indicates that the flaky and transparent substance is B2O3 due to EDS's low sensitivity to light elements of boron. The surface of samples oxidized at 1000 °C for 30 min appears rugged, and many pores are observed and a smooth layer is not formed (Fig. 5(c) and (d)). EDS analysis of the sample surface evidences the presence of ZrO2 grains, and some of them embeds in a glassy phase, which contained oxygen, silicon and boron being not detectable by the apparatus (Fig. 5(d)). As the temperature approaches 1200 °C, the pores significantly reduces and some of them are filled with liquid phase. When the temperature further rises to 1300 °C, the pores are almost filled with glass phase, a dense layer forms and completely covers the whole surface. XRD patterns of the surface of ZrC + SiC + ZrB2 composite oxidized in the 800 °C-1300 °C temperature range for 30 min are shown in Fig. 6. At 800 °C, ZrO2 peaks appear along with ZrC and SiC ones. However, B2O3 phase is not detected by XRD analyses on the exposed surface, which is consistent with the result observed by Gao et al. [23]. It is believed that the samples are quenched in air during the course of drawing the samples rapidly out of the furnace and that the liquid B2O3 is not able to crystallize before solidification. ZrC peaks are no more visible over 1000 °C. SiC peaks are seemingly visible at 1000 °C and disappear higher than this temperature. However, it should be noted that SiC  particles are not clearly observed in SEM microstructure at 1000 °C. This may be because SiC particles remain embedded in the ZrO2-glass scale. The intensity of the diffraction peaks of ZrO2 phase gradually increases with increasing temperature and, of the detected phases, is the strongest. When the samples are exposed to air at or above 1200 °C, there are only diffraction peaks of ZrO2 phase on the oxidation surface. The borosilicate is not detected since it is amorphous [24]. In order to deeply understand the oxidation mechanism of ZrC ceramic composites, it is of great interest for the observation and analysis of the structure of oxide layer and the morphology of each layer. Fig. 7 shows the cross-sectional micrographs of ZrC + SiC and ZrC + SiC + ZrB2 composites oxidized at 800 °C for 30 min. It is observed that the oxide layer thickness is ~ 265 μm for ZrC + SiC composite and ~ 240 μm for ZrC + SiC + ZrB2 composite after oxidation at 800 °C for 30 min, respectively. The cross-sectional structures of oxide layer in both the composites are not dense, and some microcracks and voids are visible. However, comparatively, the cross section of ZrC + SiC + ZrB2 composite seems more loose, which results in low strength and high shedding trend of oxide layer. The reason for relatively thinner oxide layer of ZrC + SiC + ZrB2 composite is presumably the structure of oxide layer. According to the EDS analysis, the oxide scale of ZrC + SiC composite only consists of ZrO2 layer with unoxidized SiC particles; however, ZrC + SiC + ZrB2 composite is covered by ZrO2 layer filled with B2O3, in which embeds unoxidized SiC particles. The presence of B2O3 acts as an effective barrier against oxygen diffusion to a certain extent [15,25]. Fig. 8 show s the t ime dependen ce o f ox ide sca le th i ckne s s o f Z rC + S iC + Z rB 2 compo s i te a t 1200 °C and 1300 °C . Obv iou s ly , the variation of oxide scale thickness implies that the growth rate of oxide scale markedly decreases with longer exposure time and higher exposure temperature. This means that the oxidation resistance of ZrC + SiC + ZrB2 composite is better under static oxidation conditions with high temperature. The cross-sectional morphology of ZrC + SiC + ZrB2 composite oxidized at 1200 °C for 30 min is shown in Fig. 9. The combination of SEM and EDS of total oxidized cross section reveals the composition of each region at this temperature as follows: A is the dense compact ZrO2-glass assembly; B is the ZrO2-glass assembly contained with large amounts of B2O3 flakes; C is the flaky B2O3 embedded in ZrO2-glass assembly; D and E are similar to B and C; and F is the transition layer, viz. being oxidized layer. EDS shows that Zr, O and Si are present as the primary elements in the oxidized layer, as shown in Fig. 9(b). Although B is not be detected by EDS due to the limit of the detection capability, a certain amount of B element may still exist as borosilicate. Based on above observation and analysis, it can be concluded that the oxide scale is composed of (1) ZrO2-glass assembly, (2) ZrO2-glass assembly with B2O3 flakes, (3) being oxidized layer and (4) unaffected ZrC + SiC + ZrB2 composite.  Fig. 4. High magnification SEM micrographs of the surfaces for ZrC + SiC composite after oxidation for 30 min at (a) 1200 °C and (b) 1300 °C.  \\x0c\", 'B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  163  (a)  (c)  (e)  (g)  (b)  (d)  (f)  (h)  Fig. 5. SEM micrographs of the surfaces for ZrC + SiC + ZrB2 composite after oxidation for 30 min at (a and b) 800 °C, (c and d) 1000 °C, (e and f) 1200 °C, and (g and h) 1300 °C.  The cross-sectional structure of ZrC + SiC + ZrB2 composite at 1300 ºC for 30 min is similar to that of oxidation at 1200 ºC. The difference is only that the number of flaky B2O3 reduces and ZrO2 particles grow, as shown in Fig. 10.  Oxidation mechanism  The oxidation resistance of the material is closely related to the types and characteristics of oxide formed under aerobic conditions. The  \\x0c', '164  B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  Fig. 8. Oxide scale thickness of ZrC + SiC + ZrB2 composite with respect to the oxidation time at 1200 °C and 1300 °C.  Fig. 6. XRD patterns of the surfaces of ZrC + SiC + ZrB2 samples oxidized at 800 °C- 1300 °C after 30 min.  chemical reaction taking place during the oxidation process is essential because it is of value to help to understand and analyze the oxidation mechanisms of materials systems. Therefore, it is necessary to study the thermodynamics of the system to determine the possibility of some reactions during the oxidation process. For ZrC + SiC and ZrC + SiC + ZrB2 materials, the possible reactions of the components describing the oxidation process are given as follows:  SiC sð  Þ þ O2 gð  Þ→SiO2  sð  Þ þ CO2 gð  Þ  SiC sð  Þ þ O2 gð  Þ→SiO2  SiC sð  Þ þ O2 gð  Þ→SiO2  sð Þ þ CO gð Þ  sð Þ þ C sð Þ  ð4Þ  ð5Þ  ð6Þ  ZrC sð  Þ þ O2 gð  Þ→ZrO2  sð  Þ þ CO2 gð  Þ  ZrC sð  Þ þ O2 gð  Þ→ZrO2  sð Þ þ C sð Þ  sð Þ þ O2 gð  Þ→ZrO2  sð  Þ þ B2O3  lð Þ  ZrB2  ð1Þ  ð2Þ  ð3Þ  On the basis of available thermodynamic data [26], the changes in standard Gibbs free energy ΔG0 and reaction enthalpy ΔH0 for each react ion are ca lculated and shown in F ig . 11 as a funct ion o f tempe ra tu re . I t can be found tha t in the inves t iga ted tempe ra tu re range , a l l reac t ions a re the rmodynam ica l ly favo rab le (ΔG 0 b 0 ) , and are exothermic (ΔH0 b 0) . In the ZrC + SiC system, compared with SiC, the oxidation resistance of ZrC is poor, and it is preferred to react with oxygen to form ZrO2 and  Fig. 7. Low and high magnification cross-sectional micrographs of ZrC composites oxidized at 800 °C for 30 min: (a and b) ZrC + SiC, and (c and d) ZrC + SiC + ZrB2.  \\x0c', 'B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  165  Fig. 9. SEM images and EDS analysis from polished cross section of ZrC + SiC + ZrB2 composite after oxidation at 1200 °C for 30 min: (a) a general view, and (b-g) details of A-F zones in (a).  carbon oxides in static air below 1300 °C. The oxidation from ZrC to ZrO2 is accompanied by 28% volume expansion and the release of gaseous products. Moreover, ZrO2 has a high thermal expansion coefficient, and as the temperature decreases, the tetragonal-to-monoclinic phase  transformation is accompanied by 5% volume expansion [27]. Therefore, the volume expansion derived from the formation of stable ZrO2 and the upwelling of aggregate gaseous oxides result in the fracture of particles and the separation of layers, with consequent formation of cracks and  \\x0c', '166  B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  Fig. 10. SEM images of polished cross section of ZrC + SiC + ZrB2 composite after oxidation at 1300 °C for 30 min: (a) low magnification of the cross-section, and (b-d) details of A-C zones in (a).  pores (as shown in Fig. 7). With the progress of oxidation, the volume expansion which develops in the oxide layer increases and eventually leads to form a porous ZrO2 scale which allows gaseous diffusion of oxygen through the pores to the ZrC surface, and therefore provides no oxidation protection [16]. SiC begins to markedly oxidize as the temperature approaches ~ 1300 °C [28]. The oxidation products are SiO2 glass and carbon oxides. The melting point of SiO2 is 1713 °C, which means that the viscosity of SiO2 is very high and is difficult to spread in the material surface for oxide film. Additionally, compared with ZrO2, SiO2 is extremely insufficient at a temperature of below 1300 °C to form a dense tight oxidation layer. Therefore, SiO2 from SiC exerts no effective protection with increasing oxidation temperature and prolonging holding time. The dominant chemical processes of ZrC + SiC composite in static air below 1300 °C are oxidations of ZrC and SiC by reactions (1), (4) and (5).  The oxidation process of ZrC + SiC + ZrB2 composite is different from that of ZrC + SiC composite. ZrB2 begins to oxidize to ZrO2 and liquid B2O3 in air at 800 °C by reaction (3). B2O3 is known to have a low melting point of 450 °C and high vapor pressure. At temperatures below 1100 °C, the glassy B2O3 film is an effective barrier to transport of oxygen. At temperatures above 1100 °C, B2O3 rapidly vaporizes according to reaction (7), thus reducing the effectiveness of the diffusion barrier [29,30].  B2O3  lð Þ→B2O3 gð  Þ  ð7Þ  For ZrC + SiC + ZrB2 material system, in the initial oxidation stage, ZrC and ZrB2 are oxidized to form ZrO2 and B2O3 that covers the surface of ZrC + SiC + ZrB2 composite and released gaseous oxides escape the surface. However, the oxidation rate of ZrB2 by reaction (3) is much lower than that of ZrC by reaction (1) between 800 °C and 1000 °C. In  Fig. 11. Variation of (a) ΔG0 and (b) ΔH0 of different oxidation reactions as function of temperature.  \\x0c', 'B. Ma et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 159-167  167  this case, the generated liquid B2O3 is not enough and only distributes in the local region, so they cannot form continuous phase and not completely spread in the material surface to effectively prevent the oxidation. As the oxidation temperature increases to 1000 °C, the oxidation rate of ZrB2 increases. Correspondingly, both ZrC and ZrB2 are subjected to dramatic oxidation reaction at this temperature, resulting in the accelerated growth of ZrO2 and the rapid increase of volume effect. Except these, the thermal stresses are generated due to the large difference in the thermal expansion coefficients of different components during cooling. The above factors eventually induce the shedding of oxide scale. Although the borosilicate glass can be formed from the chemical reaction between SiO2 and B2O3 (reaction (8)) at 1000 °C, the formation rate of SiO2 is slow; therefore, there presents less borosilicate glass in oxide scale, which cannot form the oxide film to protect the inside structure. When the oxidation temperature is up to 1200 °C, more SiC particles oxidize, thus promoting the formation of more borosilicate glass. The formation of this glass can exert more protective effect for the surface of samples at this temperature. The provided protection mechanism could be called self-healing protection, namely, when the composite is exposed to high-temperatures, a new fresh stable layer instantly forms and heals the preexisting cracks [22]. When the oxidation temperature continues to increase to 1300 °C, B2O3 should start to vaporize more rapidly; however, the formation of a significant amount of borosilicate glass reduces the vaporization of B2O3 and exerts a protecting effect, thus lowering the oxidation rate. Simultaneously, the partial vaporization of B2O3 causes the emergence of SiO2-rich borosilicate glass that is more stable than the B2O3-rich borosilicate [22]. The combination of above phenomenon makes ZrC + SiC + ZrB2 composite obtained thinner oxide layer and less mass gain at 1300 °C. The presence of ZrB2 particle is certainly responsible for rather good oxidation resistance. It further shows that the oxidation resistance of ZrC + SiC + ZrB2 composite is higher than that of ZrC + SiC composite.  B2O3  lð Þ þ SiO2  sð  Þ→SiO2 \\x01 B2O3  lð Þ  ð8Þ  Conclusions  Hot-pressed ZrC + SiC and ZrC + SiC + ZrB2 composites were prepared and the oxidation behavior of both the materials was investigated at temperatures in the range 800 °C − 1300 °C in laboratory air. It is found that the mass gain per surface area of ZrC + SiC composite is lower than that of ZrC + SiC + ZrB2 composite at the same time at 1000 °C, showing better oxidation resistance comparatively. At 1200 °C, the mass gain of ZrC + SiC + ZrB2 composite changes according to the change of oxidation time. After oxidation for 30 min, the mass gain of ZrC + SiC + ZrB2 composite is followed by a reduction and is lower than ZrC + SiC composite. When the temperature increases to 1300 °C, the mass gain of ZrC + SiC + ZrB2 composite is significantly lower than that of ZrC + SiC composite, and its oxidation behavior obeys the parabolic law. Meanwhile, the results of oxidized variation in microstructure have revealed that ZrC + SiC + ZrB2 composite is more resistant to oxidation than ZrC + SiC composite at both 1200 °C and 1300 °C. The main oxidation products of ZrC + SiC composite are ZrO2 and SiO2, while that of ZrC + SiC + ZrB2 composite are ZrO2, B2O3, SiO2 and borosilicate glass. The oxide scale of ZrC + SiC + ZrB2 composite is more protective than that of ZrC + SiC composite because the formation of borosilicate glass seals pores and coats the sample surfaces, greatly limiting the diffusion of oxygen into the internal materials.  The introduction of ZrB2 particles markedly improves the resistance to oxidation of ZrC + SiC + ZrB2 composite at high temperature.  Acknowledgements  This work is supported by the Natural Science Heilongjiang Province, China (no. E201021).  Foundation of  References  [14]  [1] Savino R, Fumo MDS, Paterna D, Serpico M. Aerothermodynamic study of UHTC-based thermal protection systems. Aerosol Sci Technol 2005;9:151-60. [2] Talmy IG, Zaykoski JA, Opeka MM. Synthesis, processing and properties of TaC-TaB2-C ceramics. J Eur Ceram Soc 2010;30:2253-63. [3] Upadhya K, Yang JM, Hoffmann WP. Materials for ultrahigh temperature structural applications. Am Ceram Soc Bull 1997;76(12):51-6. [4] Ma BX, Han WB. Thermal shock resistance of ZrC matrix ceramics. Int J Refract Met Hard Mater 2010;28:187-90. [5] Zhao LY, Jia DC, Duan XM, Yang ZH, Zhou Y. Pressureless sintering of ZrC-based ceramics by enhancing powder sinterability. Int J Refract Met Hard Mater 2011;29:516-21. [6] Landwehr SE, Hilmas GE, Fahrenholtz WG, Talmy IG, Dipietro SG. 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},{
  "_id": 154,
  "PDF": "Oxidation behaviors of (Hf0.25Zr0.25Ta0.25Nb0.25)C and (Hf0.25Zr0.25Ta0.25Nb0.25)C-SiC at 1300-1500 degrees C.pdf",
  "Text": "['Journal of Materials Science & Technology 60 (2021) 147-155  Contents lists available at ScienceDirect  Journal of Materials Science & Technology  j o u r n a l h o m e p a g e : w w w . j m s t . o r g  Research Article  Oxidation behaviors of (Hf0.25Zr0.25Ta0.25Nb0.25)C and (Hf0.25Zr0.25Ta0.25Nb0.25)C-SiC at 1300-1500  C  Haoxuan Wang a , Shouye Wang a , Yejie Cao a,∗ , Wen Liu b , Yiguang Wang c,∗  a Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China b School of Materials Science and Engineering, Zhengzhou University, Zhengzhou, Henan 450052, PR China c Institute of Advanced Structure Technology, Beijing Institute of Technology, Haidian District 100081, Beijing, PR China  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 7 March 2020 Received in revised form 21 May 2020 Accepted 21 May 2020 Available online 7 July 2020  Keywords:  High-entropy carbide Oxidation resistance Oxidation mechanism  In this work, high-entropy ceramics (Hf0.25 Zr0.25 Ta0.25Nb0.25 )C (HZTNC) and HZTNC doped with 20 vol% SiC (HZTNC-SiC) were fabricated by spark plasma sintering. Their oxidation behavior was  investigated over the temperature range of 1300-1500   C for up to 60 min. Both HZTNC and HZTNC-SiC exhibited good oxidation resistance, and their weight change as a function of oxidation time obeyed a parabolic law. Through XRD, microstructure observation, and elemental mapping analysis of the oxide  layers,  it was found that the formation of Nb2 Zr6O17 , Hf6 Ta2O17 , and (Ta, Nb)2O5 mixed-oxide  layers effectively protected the matrix from further oxidation. Microcracks began to appear on the oxide layer of HZTNC at high temperatures after 60 min of oxidation. However, the addition of SiC in HZTNC suppressed these microcracks at high temperatures due to the active oxidation of SiC. Compared with the oxides formed on HZTNC,  the additional  formation of Hf(Zr)SiO4 on HZTNC-SiC could  further  improve  its oxidation resistance over HZTNC ceramics. © 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.  1.   Introduction  The unique properties of high-entropy alloys (HEAs) have stimulated  the  fast development of high-entropy ceramics  (HECs)  in recent years  [1-9]. HECs are single-phase compounds composed of cations or anions of  four or more  types of elements at either an equimolar or near-equimolar  ratio. The HECs,  like HEAs, are expected to exhibit  improved mechanical and  functional properties due to the “cocktail” effect [10-16]. In 2015, Sarkar et al. first developed  (MgNiZnCuCo)O anode materials with excellent storage capacity [17]. Since then, several HECs with different crystal structures have been synthesized, including materials with a rock salt  structure  [18-22], perovskite  structure  [23], fluorite  structure [24], spinel structure [25], AlB2 structure [1,26], half-Heusler structure [27], and CrSi2 structure [28]. These ceramics exhibited better mechanical properties  [29],  ionic conductivities  [30], corrosion resistance [31-35], nuclear waste immobilization [36], and low temperature conductivity [37] than the corresponding singlecomponent ceramics.  ∗ Corresponding authors. E-mail addresses: Caoyejie@nwpu.edu.cn (Y. Cao), wangyiguang@bit.edu.cn (Y. Wang).  Ultrahigh-temperature HECs (UHTHECs) have attracted particular attention due  to  their application  in areas such as  thermal protection systems of hypersonic vehicles [38], high-temperature devices [39], and radiation-resistant materials for advanced nuclear energy systems [40,41]. Ultrahigh-temperature ceramics are usually transition-metal carbides or borides, which are easily oxidized in oxygen-containing environments [42-45]. Thus, improving the oxidation resistance of these materials  is one of the key research topics  in  this field. Luo et al.  [42]  synthesized  the first boride UHTHEC, which had a significantly improved oxidation resistance compared with the corresponding single-component borides. Similarly, UHTHEC carbides also exhibited better oxidation resistance than the corresponding single-component carbides [35,42]. Since complex oxides are  formed during  the oxidation process due  to the multi-cation nature of the UHTHEC,  it can be argued that  its underlying oxidation mechanism  is controlled by oxygen  inward diffusion  [46] or  the cations outward diffusion  [47]  through  the oxide layer. However, it is generally believed that the formation of complex oxides during the oxidation of UHTHECs and the subsequent hindered diffusion effect of high-entropy cations account for the improvement in oxidation resistance [31,35]. In our previous study [35], a (Hf0.2 Zr0.2 Ti0.2 Ta0.2Nb0.2 )C (HZTTNC) UHTHEC was prepared with  improved oxidation resistance compared  to  the  individual  carbides, which was  still  inferior  https://doi.org/10.1016/j.jmst.2020.05.037 1005-0302/© 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.                                    \\x0c', '148   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  compared  to  that of UHTCs with  the addition of SiC  (such as ZrB2 -SiC) [47-49]. When SiC was added to HZTTNC material [31], the resultant ceramics (HZTTNC-SiC) showed  improved oxidation resistance, albeit still not as good as that of ZrB2 -SiC. The outward diffusion of TiO was proposed to be the dominant process  in the oxidation of both HZTTNC [35] and HZTTNC-SiC [31]. Due to this process, the addition of SiC in HZTTNC did not lead to an improvement  in the oxidation resistance of HZTTNC compared to that of ZrB2 -SiC ceramics. With the aim of  further  improving the oxidation  resistance, TiC was  removed  from HZTTNC  to  form  (HZTNC) materials  in  this work,  thus avoiding the formation of TiO during oxidation. HZTNC doped with 20 vol% SiC (HZTNC-SiC) was also prepared. The oxidation behavior of the prepared materials was assessed  in the temperature range of 1300-1500  C to evaluate whether this strategy could  further enhance the oxidation resistance of UHTHECs.  It  is expected that this research would be helpful to  improve the high-temperature properties of UHTHECs by adjusting their composition.  Zr0.25 Ta0.25Nb0.25 )C   (Hf0.25  2. Experimental methods  Mechanical alloying combined with spark plasma sintering was used to synthesize HZTNC and HZTNC-SiC. Four carbide powders of HfC, ZrC, TaC, and NbC (purchased from HWRK CHEM, Beijing, China; 99 %, average size: 2  \\u242em) were weighed at a 1:1:1:1 M ratio of HfC to ZrC to TaC to NbC. Ball-milling was performed in a planetary ball mill (PMDW, Nanjing, China) at a speed of 500 rpm for 18 h. In order to prevent the vials from overheating, the milling was stopped for 5 min every 15 min. The balling medium was tungsten carbide and alcohol was used as dispersant. After ball-milling, the obtained powder mixture was divided into two parts: one part was used for the sintering of HZTNC bulk and the other part was mixed with SiC powder (99 %, 2  \\u242em, HWRK CHEM, Beijing, China) at a volume fraction of 80:20. The powder mixed with SiC was placed in a nylon tank and ball-milled at 130 rpm for 24 h in a roller ball mill (GQM-2−5, TENCEN POWDER, Changsha, China). The as-milled powders (with and without SiC) were placed in a  Ф20 graphite die  Fig. 1. XRD patterns of the as-received ceramics.  ×  ×  and sintered by spark plasma sintering (Beijing Ailin Furnace Technology Co, Ltd.) at 2000  C for 15 min under a uniaxial pressure of 35 MPa in vacuum at a heating rate of 150  C/min. The as-sintered HZTNC and HZTNC-SiC ceramics were cut into 10 mm   5 mm   5 mm bulk samples for oxidation tests. One side of the sample was polished to 1  \\u242em finish with diamond lapping films. The polished samples were then placed on hollow zirconia balls in a zirconia crucible with the polished side facing up. Isothermal oxidation was carried out in an alumina tube furnace (GSL-1600X, Hefei Kejing Materials Technology Co., Ltd., Hefei, China) at 1300-1500  C  in air. When  the preset  temperature was reached,  it was  initially maintained for 10 min and the furnace tube was evacuated. Then, the furnace tube was quickly filled with air within 30 s, and finally adjusted with flowing air at a rate of 10 mL/min. Following the completion of oxidation, the air was replaced with high-purity argon and the samples were cooled to room temperature. The sam Fig. 2. SEM images showing the polished surface of HZTNC and HZTNC-SiC ceramics (a) HZTNC, (b) HZTNCSiC.  \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   149  Fig. 3. SEM surface images of HZTNC and HZTNC-SiC ceramics after isothermal oxidation at different temperatures: (a) HZTNC, 1300   C, (b) HZTNC-SiC, 1300   C, (c) HZTNC, 1400   C, (d) HZTNC-SiC, 1400   C, (e) HZTNC, 1500   C, (f) HZTNC-SiC, 1500   C.  Table 1 Calculated parabolic rate constants and activation energies for HZTNC, HZTNC-SiC, HZTTNC [35], HZTTNC-SiC [31] and ZrB2 -SiC [46] ceramics.  kp, 1300   C  (mg2 /cm4 h)   kp, 1400   C  (mg2 /cm4 h)   kp, 1500   C  (mg2 /cm4 h)   Activation energy (kJ/mol)  HZTNC  HZTNC-SiC  HZTTNC  HZTTNC-SiC  ZrB2 -SiC   26.31  6.36  53.90  18.30  11.81   47.46  15.03  81.76  32.67  -   64.07  22.42  152.10  56.69  38.91   103 146 134 130 212  ples were then weighed by an analytical balance with an accuracy of 0.01 mg (Mettler Tolendo AG135, Greifensee, Switzerland). The above oxidation procedure was repeated five  times  for  the  total oxidation time of 60 min at 1300  C, 1400  C, and 1500  C, respectively. All oxidized samples were embedded in epoxy resin and cut in the middle. The cross-section was then polished to 1  \\u242em finish with diamond lapping films. The HZTNC and HZTNC-SiC phases after  isothermal oxidation were characterized by X-ray diffraction (XRD; Rigaku D/max-2400, Tokyo, Japan) with Cu Ka radiation. The microstructure of the samples was characterized by scanning electron microscopy (SEM; JEOL 6700 F, Tokyo, Japan) together with energy dispersive spectroscopy (EDS) for elemental mapping.  3. Results and discussion  ±  ±  By using the Archimedes method, the relative densities of the HZTNC and HZTNC-SiC bulks were measured to be 9.85   0.27 (>98 % theoretical density) and 8.59   0.27 (>99 % theoretical density), respectively. The XRD patterns of HZTNC and HZTNC-SiC showed only one phase with FCC structure for HZTNC, similar to that of TaC, albeit with a slight shift of the diffraction peaks toward the right (Fig. 1). Compared to HZTTNC in our previous study [35], despite the removal of TiC, the remaining four components retained the singlephase HEC structure. The sharp peaks observed for HZTNC-SiC were assigned to HZTNC and the weak peaks were attributed to SiC; no other peaks were observed  in HZTNC-SiC (Fig. 1). The morphologies of the polished HZTNC and HZTNC-SiC surfaces (Fig. 2) were also characterized. The HZTNC ceramic showed uniform elemental distribution (Fig. 2(a)), while the HZTNC-SiC ceramic showed two regions, a gray region corresponding to the HZTNC phase and a dark region corresponding to the SiC phase (Fig. 2(b)). A few micropores were observed on the surface of HZTNC, whilst HZTNC-SiC was almost fully dense,  indicating that the addition of SiC helped to increase the relative density of the ceramics. Fig. 3 shows the morphologies of HZTNC and HZTNC-SiC samples after isothermal oxidation at 1300−1500  C for 60 min. It can be seen that the formed oxide layers were dense without obvious for both HZTNC and HZTNC-SiC oxidized at 1300  C and defects   \\x0c', '150   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 4.  (a-c) Square of the specific weight change as a  temperature.  function of oxidation time at 1300, 1400, and 1500   C;   (d) the relationship between   lnkp and the reciprocal of  1400  C. At 1500  C, a number of microcracks were observed on the surface of oxidized HZTNC, while the surface of HZTNC-SiC was still dense without obvious defects. The isothermal oxidation data of HZTNC and HZTNC-SiC at 1300  C, 1400  C, and 1500  C were plotted in the form of the square of the specific weight change as a function of oxidation time (Fig. 4). These data were fitted based on the following equation [50]:  w2=   kp t   (1)  t   where w  is  the specific weight change during oxidation,  is  the oxidation time, and kp is the parabolic oxidation rate. As seen from Fig. 4, these curves were  linearly fitted, showing that the oxidation process followed the parabolic rate law (Eq. (1)). The parabolic oxidation behavior indicates that the diffusion process dominated the oxidation of HZTNC and HZTNC-SiC. The kp values were calculated through the slope of the w2 -t plots and compared with the kp values for HZTTNC and HZTTNC-SiC [31,35] (Table 1).  Indeed, the for HZTNC at 1300-1500  C were  kp values  lower than those  for HZTTNC, indicating that the removal of TiC led to an improvement in the oxidation resistance of UHTHECs. HZTNC-SiC also showed better oxidation resistance than HZTTNC-SiC, and even superior to that of ZrB2 -SiC ceramics [44-49]. The  increase  in kp value with  the  increase  in oxidation  temperature  indicates  the  faster  oxidation  of  ceramics  at  higher temperatures. The temperature dependence of kp can be described by Arrhenius equation, as follows [50]:  kp=  k 0 exp(-E /RT   )   (2)  where k0 is a pre-exponential constant, E is the oxidation activation energy, R  is the gas constant, and T  is the oxidation temperature. The natural logarithm of kp at 1300-1500  C (ln kp ) was plotted as a function of reciprocal temperature (1/T) in Fig. 4(d). The oxidation  activation energies of HZTNC and HZTNC-SiC were calculated from the slopes of the curves in Fig. 4(d) and listed in Table 1. The oxidation activation energies were 103 kJ/mol for HZTNC, 146 kJ/mol for HZTNC-SiC, and 134 kJ/mol for HZTTNC [31,35]. The TiO outward diffusion during the oxidation of HZTTNC [35] cannot be the dominant step  in the oxidation mechanism of HZTNC and HZTNC-SiC due to the absence of TiC in both ceramics. Additionally, given their different oxidation activation energies, the controlling step in the oxidation of HZTNC would also be different from that in HZTNC-SiC. The outer oxide layers of HZTNC and HZTNC-SiC ceramics oxidized at 1300-1500  C  for 60 min were characterized by XRD (Fig. 5). The main phases of the outer oxide  layer of HZTNC were Nb2 Zr6O17 and Hf6 Ta2O17 , with a  small amount of  (Nb,Ta)2O5 (Fig. 5(a)). Without  the outward diffusion  and  evaporation of TiO,  like  in HZTTNC,  the outer  layer of Nb2 Zr6O17 , Hf6 Ta2O17 , and  (Nb,Ta)2O5 oxides mixture was dense, providing  improved oxidation protection  for HZTNC compared  to HZTTNC. With  the increase in oxidation temperature, the concentration of (Nb,Ta)2O5 (Fig. 5(a)) was  reduced,  indicating  that  the  formation  rate of (Nb,Ta)2O5 was inhibited at higher temperatures. Thus, Ta and Nb tended to form Nb2 Zr6O17 and Hf6 Ta2O17 .  In addition to the oxidation products of HZTNC  in the oxidized HZTNC-SiC samples, a new phase of (Hf,Zr)SiO4 was also observed (Fig. 5(b)). The solid solution (Hf,Zr)SiO4 was  likely derived  from the reactions of the oxidation products of HfC, ZrC, and SiC in HZTNC-SiC. The presence of (Hf,Zr)SiO4 in the oxide layer enhanced the oxidation resistance of HZTNC-SiC. The cross-sectional morphologies and  the corresponding element mapping of the HZTNC and HZTNC-SiC ceramics oxidized at 1300  C  for 20 min and 60 min were analyzed (Figs. 6 and 7). A dense oxide layer was formed on the top of HZTNC or HZTNC-SiC ceramics after both 20 min and 60 min. Based on the elemental    \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   151  Fig. 5. XRD patterns of oxidized HZTNC and HZTNC-SiC ceramics at different temperatures: (a) HZTNC, (b) HZTNC-SiC.  Fig. 6. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1300   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  \\x0c', '152   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 7. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNCSiC ceramics at 1300   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min.  Fig. 8. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1400   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  mapping analysis, Hf, Zr, Ta, and Nb were found to be distributed uniformly in the oxide layer of HZTNC ceramics. Thus, according to the XRD results, it was hypothesized that HfO2 reacts with Ta2O5 to form Hf6 Ta2O17 , and ZrO2 reacts with Nb2O5 to form Nb2 Zr6O17 in the oxide  layer. Subsequently, the unreacted Ta2O5 and Nb2O5 form the solid solution of Nb(Ta)2O5 . As seen from the elemental mapping results of the oxidized HZTNC-SiC ceramics, the elements of Zr, Hf, and Si were enriched in the oxide layer while Ta and Nb  were also detected  in the oxide  layer. This result  is  in agreement with the XRD analysis which showed that the oxide phases were (Hf,Zr)SiO4 , Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 . At 1400  C, the oxide layers on HZTNC and HZTNC-SiC increased in thickness compared to those at 1300  C (Figs. 8 and 9), whereas the elemental distribution and oxide structure were similar to those at 1300  C. The cross-sectional morphologies of  the HZTNC and HZTNC C were  samples  oxidized  at  1500 observed  by  SEM  SiC   \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   153  Fig. 9. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNCSiC ceramics at 1400   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min.  Fig. 10. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1500   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  (Figs. 10 and 11). The oxide  layer of HZTNC at 20 min was dense with uniformly distributed elements. Following 60 min of oxidation time, two oxide  layers were observed on the HZTNC surface: a Zrand Hf-enriched top layer with some lateral microcracks and pores, and a relatively dense second layer formed by a mixture of Hf, Zr, Ta, and Nb oxides. Due to the absence of TiO diffusion during oxidation, the formed oxide layer of HZTNC was dense after the first 20 min. With the increase in oxidation time at 1500  C, the thick ness of  the dense  layer  increased. Generally, stress  is developed with the growth in the oxide layer due to the difference in properties of the matrix and the oxides on  its top [51-53]. It  is believed that the stress in the dense oxide layer of HZTNC increases with the increase  in thickness, finally becoming  large enough to generate microcracks in the oxide layer. These microcracks provide diffusion channels for oxygen to further oxidize HZTNC to form a dense second oxide layer. For HZTNC-SiC ceramics, the oxide layers formed  \\x0c', '154   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 11. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC-SiC ceramics at 1500   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min. (*1 Hf(Zr)TiO4 , Hf(Zr)SiO4 , Nb(Ta)2 O5 layer; 2 Hf-Zr-Ta-Nb-Si-O layer).  Fig. 12. Schematic oxide structure of HZTNC and HZTNC-SiC ceramics: (a) HZTNC, (b) HZTNC-SiC.  at 20 min and 60 min were both dense without obvious cracks. When the ceramic was oxidized for 60 min, two oxide layers were observed. The outermost dense layer was mainly composed of Zr, Hf, and Si, corresponding to (Hf,Zr)SiO4 solid solution according to XRD analysis. The second layer consisted primarily of Hf, Zr, Ta, and Nb oxides based on elemental analysis, and was depleted in the Si element. Similar phenomena were also observed  in the oxidation of ZrB2 -SiC ceramics [44,46-49,54], in which the Si-depleted layer was detected. Therefore, it is proposed that SiO outward diffusion occurred during the oxidation of HZTNC-SiC ceramics. The outward diffusion and evaporation of SiO may release the growth stress and avoid the formation of microcracks as the oxide layer thickens. Based on the above discussion, the plausible oxidation mechanisms for HZTNC and HZTNC-SiC are proposed, as shown in Fig. 12. According to the oxidation kinetics, the diffusion process should be dominant during the oxidation of HZTNC. It is unclear whether the controlling process is the inward oxygen diffusion or the outward  diffusion of cations such as Hf/Zr. However, the oxidation activation energy of HZTNC is only 103 kJ/mol, much lower than that required for cation outward diffusion [35]. Therefore, the  inward diffusion of oxygen through the  formed oxide  layer  is  likely the dominant process for the oxidation of HZTNC (Fig. 12(a)). During the oxidation process, a dense complex oxide layer composed of Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 is  formed on HZTNC, preventing  the matrix  from  further oxidation. As the dense  layer grows, growth stress accumulates and becomes  large enough to break the  layer and  form microcracks. A new, dense oxide  layer  is  then  formed below the  former  layer by the diffusion of oxygen through these cracks for further oxidation. The formation of the Si-depleted layer indicates that the oxidation mechanism of HZTNC-SiC is similar to that of ZrB2 -SiC, wherein the outward diffusion of SiO  is the controlling step during oxidation [44,54]. At low temperatures, the diffusion of SiO is so slow that the Si-depleted layer is not obvious and only a dense layer can be  \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   155  observed. At high temperatures and long oxidation times, the diffusion of SiO is accelerated and the Si-depleted layer becomes more obvious. The diffused SiO reacts with the Zr/Hf oxides to form an additional dense  layer of (Hf, Zr)SiO4 , thus preventing the matrix from oxidation. Moreover, the outward diffusion of SiO also releases oxide growth stress, ensuring the integrity of the protective oxide layer. The oxidation resistance of HZTNC-SiC is thereby enhanced compared to that of ZrB2 -SiC ceramics.  4. Conclusion  In conclusion, the oxidation behaviors of HZTNC and HZTNC-SiC ceramics were assessed over the temperature range of 1300-1500  C in air for 1 h. The results indicated that both HZTNC and HZTNCSiC ceramics exhibited parabolic oxidation behavior over the tested temperature range. The inward diffusion of oxygen was proposed as  the possible oxidation mechanism  for HZTNC ceramics,  leading to the  formation of a dense mixed oxide  layer of Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 . The accumulated growth stress could break the oxide layer at high temperatures and long oxidation time, resulting in cracks and further oxidation of the matrix below. After the addition of SiC in HZTNC, the outward diffusion of SiO became the controlling process. This further improved the oxidation resistance of the ceramics and released the growth stress, thus avoiding the generation of microcracks. The multiple components of HECs can synergistically affect the properties of the materials. The present study provides an example of how the oxidation resistance of UHTHECs can be  improved by removing TiC  from HZTTNC to  form a new HEC, thus demonstrating that the properties of HECs can be adjusted by tuning their composition.  Acknowledgements  This work was financially supported by the National Natural Science Foundation of China (Grant No. 51972027). We would like to thank the Analytical & Testing Center of Northwestern Polytechnical University for assistance with SEM observations.  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  "_id": 155,
  "PDF": "Oxidation behaviors of ZrB 2 –SiC binary composites above 2000 °C.pdf",
  "Text": "['Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Oxidation behaviors of ZrB2-SiC binary composites above 2000 °C  Ryo Inouea,⁎ , Yutaro Araia,b, Yuki Kubotaa,c  a Department of Materials Science and Technology, Tokyo University of Science, 6-3-1, Nijyuku, Katsushika-ku, Tokyo 125-8585, Japan b Department of Advanced Interdisciplinary studies, Graduate School of Engineering, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904,  Japan c Japan Aerospace Exploration Agency, Structures and Advanced Composite Research Unit, Ohsawa 6-13-1, Mitaka-shi, Tokyo 181-0015, Japan  A R T I C L E  I N F O  Keywords:  Ultra-high temperature ceramics (UHTCs)  ZrB2 Electric heating  Oxidation  SiC  A B S T R A C T  Rapid oxidation testing for monolithic ZrB2 and ZrB2-SiC binary composites with different SiC contents (0- 30 vol%) was performed using an electric heating system above 2000 °C. The system used in this study achieved  the high heating rate of 250 °C/s. The experimental  results  showed that  the morphologies of  the oxidized  specimens depended strongly on the SiC content. The formation mechanism of SiC-depleted layers beneath the  surface scale above 2000 °C differed completely from that below 2000 °C. Although the holding time was below  10 s, SiC-depleted layers were formed because the oxygen partial pressure of the air atmosphere was not enough to form SiO2 by the oxidation of SiC. It was determined that ZrB2-20 vol% SiC showed the best oxidation resistance above 2000 °C at high heating rates.  1.  Introduction  ZrB2-SiC binary composites (hereafter denoted as ZS) are attractive candidates for use in high-temperature applications, such as  engine  thrusters  in satellites  and  the  leading  edges  of  supersonic  vehicles. For the abovementioned applications,  the oxidation of ultra high-temperature  ceramics  (hereafter, UHTCs)  is a critical problem  because  oxidation  causes  significant material  consumption  in  the  formation  of  porous  structures.  The  oxidation mechanisms  of  ZS  composites and suitable SiC contents have been researched carefully  because excessive SiC levels cause the recession of ZS composites [1,2]. the maximum temperatures of 1200-  Isothermal oxidation tests at  1800 °C have been performed using electric furnaces in dry air, CO, and CO2 inert gas [1-6]. Previous studies showed that the oxidation of SiC began at 800 °C, forming SiO2 scale on the surface of the ZS composite. ZS maintained oxidation resistance because of  the SiO2 temperatures above 1700 °C [2]. However, Fahrenholtz  scale even at  et  al.  [7]  revealed  that  the  preferential  oxidation of  SiC occurred  beneath  the  protective  oxide  layer. They named  this  layer, which  comprised ZrB2 particles and voids evolved by the preferential oxidation of SiC, the SiC-depleted region. Thus, the oxidation of ZS  composites strongly depends on the test partial pressure of the atmosphere [1-7].  temperature and the oxygen  In actual high-temperature application conditions, the temperature  increase  rates  during  operation  are much  higher  than  those in 10-  oxidation  tests  of UHTCs  using  electric  furnaces  (typically,  100 °C/min). To reproduce this environment, arc-jet tunnel testing [8-11] and CH4-O2 torch testing [12-17] have also been performed for both monolithic ZrB2 and ZS composites. Torch tests are relatively simple and easily performed at the laboratory scale; however,  they  induce local heating in the specimen, and the formation of temperature  gradients is unavoidable. Arc-jet tunnel testing is suitable to reproduce  the re-entry environment experienced by aerospace vehicles; however,  the method requires the preparation of  large, complex-shaped speci mens.  In  addition,  the  applied  dynamic  pressure  and  oxidation  progress both strongly affect the specimen shape [8].  Recently, Karlsdottir et al. [18] developed a method for measuring  rapid  oxidation using a resistance system and a specially designed “ribbon specimen.” They oxidized ZrB2-15 vol% SiC at 1600-1700 °C for 15 min in air. They found that oxide scales compris heating  ing SiO2 and particulate ZrO2, a thin SiO2 layer, and a SiC-depleted layer were formed on the ZS composite after oxidation. Li and Chen the surface of a ZrB2-SiC-BN specimen after oxidation using a fast heating system at 1800 °C contained mainly porous ZrO2. The thickness of the oxide scale was ~200-300 µm [19]; that of the SiO2+particulate ZrO2, ZrO2+thin SiO2, and SiC-depleted layers were ~10 µm, ~100 µm, and ~100-150 µm, respectively [18]. In  also showed [19]  that  the tests by Karlsdottir et al. and Li and Chen, the initial heating rates were ~480 °C/min and ~60-80 °C/min, respectively [18,19].  Thus, depending  on heating  rate  and holding  temperature,  the  structure of the oxidized region of the ZS composite varies significantly.  To evaluate the effect of SiC content on the oxidation behaviors of ZS  http://dx.doi.org/10.1016/j.ceramint.2017.03.129  Received 11 January 2017; Received in revised form 7 March 2017; Accepted 21 March 2017  ⁎ Corresponding author.  E-mail address: inoue.ryo@rs.tus.ac.jp (R. Inoue).  Ceramics International 43 (2017) 8081-8088  Available online 23 March 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  MARK  \\x0c', 'composites, a resistance heating system is an appropriate apparatus  because it provides a heating rate in the same range of torches and arc  jet  facilities, while  avoiding  the  effect  of dynamic pressure  on the  oxidation of  the material. The  resistance heating  system may be  a  useful tool for examining the oxidation of UHTCs and reproducing high  heating rates and temperatures (above 2000 °C) in applications such as  nose  cones  or  the  leading  edges  of hypersonic  vehicles. However,  previous  studies have  only  reported behaviors  and analyses  below  1800 °C.  The objective of this study is to evaluate the oxidation resistance of  ZS  composites with  different  SiC contents  at  temperatures  above  2000 °C  under  high  heating  rates  and without  dynamic  pressure  effects. Firstly, a rapid heating system using resistance heating was  established. The developed system was applied to examine the oxida tion mechanisms of ZS composites with different SiC contents above  2000 °C at  the high heating  rate of ~250 °C/s. The  effects of high  heating rates and maximum surface  temperatures on the oxidation  mechanism of ZS composites with different SiC contents were exam ined.  2. Experimental procedure  2.1. Sample preparation  ZS binary composites with different SiC contents were fabricated by  spark plasma sintering (SPS-1030, Sumitomo Coal Mining Co. Ltd.,  Tokyo). As raw materials, commercially available ZrB2 (average particle size of ~2.1 µm, Grade F, Japan New Metals Corp., Japan) and α-SiC  (average particle size of ~22 µm, GC800, Showa Denko, Tokyo, Japan)  particles were  used.  The  compositions  of  the  ZS  composites  are  summarized  in  Table  1.  Sintering was  performed  at  2000 °C for  1800 s under Ar gas atmosphere. Disk-shaped specimens with dia meters of 15 mm and thicknesses of 3 mm were  fabricated. Fig. 1  shows  the morphologies of  the monolithic ZrB2 and ZS composites used in this study. For monolithic ZrB2, some pores exist in the ZrB2 particles. For the ZS composites, SiC particles exist in the boundaries  between  ZrB2 5 mm×5 mm×3 mmwere  particles.  Rectangular  prism specimens measuring  cut  from the as-sintered compacts using a  diamond-impregnated  blade.  The  specimens were  polished  using  diamond paste with particles of up to 1 µm. The polished surfaces  and  cross-sections  were  observed  using  an  optical microscope  (Axiovision, Carl Zeiss, Jena, Germany) and a field-emission scanning  electron microscope  (FE-SEM,  S-4200, Hitachi High-technological  Corp., Tokyo, Japan) with an acceleration voltage of 15.0 kV. The relative density, ρR, of all ZS composites was more than 95%. Here, the bulk density, ρB, is divided by the ideal density, ρI, to obtain the relative ρR. The X-ray density, diffraction profiles (not shown) from each specimen confirmed that  the  composites were not  oxidized during  processing. The average grain sizes of the ZrB2 after sintering were 3.0, 3.5, 3.4, and 3.2 µm for the ZrB2, ZrB2-10 vol% SiC, ZrB2-20 vol% SiC, and ZrB2-30 vol% SiC specimens, respectively (hereafter denoted as ZrB2, ZS10, ZS20, and ZS30). Notably, the grain size of ZrB2 after sintering was slightly increased compared to that of the raw material.  This  is  reasonable because  the SiC particles dispersed at  the grain  boundaries prevented the grain coarsening of ZrB2,  as  reported by  Zhang et al. [20].  2.2. Oxidation tests by resistance heating  The  electric  resistivity  of  the  ZS  composite  (~2-3×10−7 Ω m  [21,22])  is  low, equivalent  to that of  iron.  It  is difficult  to heat ZS  composites above 2000 °C by resistance heating because a very large  current of ~200 A is required. Fig. 2 shows a schematic of the experimental system used in this study. Two Pb-acid storage batteries  were used as power sources for  resistance heating and connected in  series,  providing  an  electromotive  force  of  ~25 V.  A monolithic  reaction-bonded silicon carbide  (RBSC)  ceramic plate was used as  the heat source to oxidize the ZS composites. The electric resistivity of RBSC is 10−1-10−3 Ω m and it is expected to reach above 2000 °C. The  temperature increase rate, maximum temperature, and holding time at  the maximum temperature were set to ~300 °C/s, ~2000 °C, and ~10 s,  respectively. The surface temperature during the test was measured  using a radiation thermometer  (measured wavelength: 0.96 µm,  IR AHS0, CHINO, Tokyo, Japan). The time-temperature relation was also  recorded using an analog data collection system (NR-500, Keyence,  Osaka, Japan). The focus spot size on the specimen surface was set as  ~5 mm. The  emissivity was  tentatively  set  at 1, because  the  value  during oxidative reactions was initially unclear. Thus, the temperature  measured by  the  radiation thermometer was  lower  than the actual  specimen temperature. We ensured that the surface temperature of the  specimen measured by  the present  system reached above 2000 °C  because the maximum temperatures of all tests exceeded 2000 °C. The  relationship  between  temperature  and  emittance  is  discussed  in  Section 3.1.  After the test,  the specimen was embedded in epoxy resin and cut  into two pieces  to observe the cross-section of  the specimen. Cross sections  of  the  post-test  specimens were  observed  using FE-SEM.  Elemental  analysis was  also  performed by  energy-dispersive X-ray  spectroscopy (EDS, AMETEC Inc., Mahwah, NJ, USA). The thickness  of  the  oxidized  layer was measured  by  image-processing  software  (ImageJ, NIH, Washington, D.C., USA). The oxidation products were  analyzed using a confocal micro-Raman spectroscopic system (NRS 7100,  JASCO Corp., Tokyo,  Japan). A blue  laser  (457 nm, Cobolt  Twist™ 50, Sweden) was used as the excitation source. The laser spot  size was ~2 µm and the exposure time of 10 s was used. The acquired  spectrum was integrated twice. The wavenumber present system was ~ ± 1.5 cm−1.  resolution in the  3. Experimental results  3.1. Surface temperature of ZS composites oxidized using resistance  heating system  Fig. 3  shows  the  typical  time-temperature  relations during  the  oxidation  tests.  The  surface  temperature  increases  rapidly when  heating begins,  reaching the maximum temperature above ~2000 °C  for ~10 s. The temperature increase rate reaches ~250 °C/s by fixing the DC voltage to ~25 V. This rate is 25-30 times higher than that in  previous studies [18]. The temperature history is independent of  the  specimen SiC contents. The actual emissivity values for the visible and infrared (IR) (wavelength: ~0.6-40 µm) of ZrB2-SiC at 1100-1500 °C in vacuum (10 Pa) are 0.7-0.8 [23]; however, these values change  depending on the temperature and on the surface oxides formed during  oxidation. For example, the emissivity values for IR (wavelength: ~1 µm) of ZrB2-SiC oxidized at 1600 °C for 10 min is 0.9-0.95 [24]. In the present research, the emissivity of the radiation thermometer is  set as 1 because  the  temperature measured by  the  thermometer  is  lower  than the  specimen temperature. As noted in Section 2.2, we  ensured that  the surface temperatures of  the specimens measured in  the present system exceeded 2000 °C by reaching maximum tempera tures exceeding 2000 °C in all tests. The relationship between emissiv Table 1  Composition and selected properties of ZS composites used in this study.  SiC content, (vol  %)  Relative density ρR (%)  Ideal density, ρI (g/ cm3)  ZrB2  0  95.6  6.09  ZrB2-10SiC(ZS10) ZrB2-20SiC(ZS20) ZrB2-30SiC(ZS30)  10  99.4  5.8  20  99.6  5.52  30  99.6  5.23  R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  8082  \\x0c', \"R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  Fig. 1. Morphologies of as-sintered monolithic ZrB2 and ZS composites (white and black arrows indicate pores and SiC particles, respectively).  2500  2000  1500  1000  ZrB2 ZS10 ZS20 ZS30  )  C  o  (  e  r  u  t  a  r  e p  m  e  T  500  0  10  20 Time (s)  30  40  50  Fig. 3. Relationship between surface temperature and time during heating. The accuracy  of the temperature is ~100 °C.  ⎛ ⎝⎜  = ε′ ε  T T'  1 / 4  ⎞ ⎠⎟  (3)  The emissivity of  the oxidized ZS composites varies from ~0.7 to  0.95. From Eq.  (3),  if  the emissivity of  the ZS composites is 0.7,  the  temperature  reaches ~2150 °C. Thus,  the minimum and maximum  temperatures reached in the present study are ~2000 and ~2150 °C,  (1)  respectively.  3.2. Morphologies and elemental analysis  Backscattered  scanning  electron  (BSE)  images  of  the  polished  cross-sections of  the ZrB2 specimen and ZS composites with different SiC contents are shown in Fig. 4(a), (b), (c), and (d), respectively.  After the oxidation test, surface scale is formed on the specimen surface  of monolithic ZrB2. Notably, the grain size of ZrB2 beneath the surface scale is much larger than that of the as-fabricated specimen. This is  reasonable because the sintering of ZrB2 begins in the temperature range of ~2000-2150 °C [20,25]. It is difficult to measure the surface  temperature directly; however, the phenomenon of sintering confirms  that the surface temperature exceeds 2000 °C.  8083  indicate the temperature, heating rate, Stefan- (5.67×10−8 W/m2 K4  and emissivity,  [24]),  re spectively. Assuming that the heating rate of the oxidation test in the present study is constant and that the emissivity is varied from ε to ε′, the temperature (T ′) measured by ε′  is obtained using Eq. (2):  ⎛ ⎝⎜  T'= qσ ε'  1/4  ⎞ ⎠⎟  (2)  From Eqs.  (1) and (2),  the relationship between temperature and  emissivity is given by:  Fig. 2. Schematic of electric heating system used in this study.  ity and temperature is expressed by the following equation:  ⎛ ⎝⎜  T= qσ ε  1/4  ⎞ ⎠⎟  where T, q, σ , and ε  Boltzmann constant  \\x0c\", 'R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  Fig. 4. Cross-sectional view of ZS composites after heating at 200 °C for 10 s. (a) Monolithic ZrB2, (b) ZS10, (c) ZS20, and (d) ZS30.  Fig. 5. Elemental mapping of Zr, Si, O, and C of (a) monolithic ZrB2, (b) ZS10, (c) ZS20, and (d) ZS30.  8084  \\x0c', 'R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  Fig. 6. Typical Raman spectra collected from grains of oxidized layers in the specimens.  depleted layer at different SiC contents.  Fig. 7. Relationships between average  grain size  of ZrB2  and porosity within SiC The morphologies of  the oxidized surfaces of  the ZS composites  differ from that of monolithic ZrB2. The ZS composites form layered structures comprising surface scales and porous subsurface layers. The  porous subsurface layer contains grains that are large compared with  those in the unreacted region. The grains become interconnected and  extensive pores are nucleated between large particles. It seems that the  grain size and porosity of  the subsurface  layer depends on the SiC  content.  Fig. 5 shows the typical EDS mapping for the post-test specimens.  The oxidized surface scale on the monolithic ZrB2 mainly contains Zr and O. Fig. 6 shows the typical Raman spectrum of a grain within the  oxidized surface scale of the monolithic ZrB2 and ZS composites. Strong peaks around 190, 305, 348, and 476 cm−1 are observed, which  completely correspond to monoclinic crystalline ZrO2 [26]. This result strongly suggests that the surface of the entire specimen is covered with  a porous ZrO2 layer. The morphologies and compositions of tested monolithic ZrB2 specimen in this study correspond to those in a report using a previous oxidation model above 1800 °C [27]. Thus, the  the post surface scale is  identified as monoclinic ZrO2; hereafter, scale is denoted as a porous ZrO2 layer. ZrO2 scale with small pores is observed in the surface scale of each ZS composite, as with the bulk ZrB2 specimen (Fig. 5(a)). The porous subsurface layers also contain Zr, while signals from additional i.e., Si, C, and O, are not detected (Fig. 5(b)-(d)). These  the surface  elements,  results suggest that the ZrB2 particles remain after oxidation and that the porous layers are formed by the oxidation of SiC particles dispersed  at the grain boundaries. In the present study, the ZS composites do not  form a protective layer of SiO2-filled ZrB2, as reported previously [7]. Hereafter, the porous subsurface layers are defined as “SiC-depleted layers.” The formation of  the SiC-depleted layer depends on the SiC  content [7].  Fig. 7 shows the size distribution of  the ZrB2 particles in the SiCthe ZS specimens. For ZS10, Figs. 4 and 5 both  depleted regions of  clearly  show that ZrO2 unoxidized region. SiO2 and ZrO2+SiO2 layers, typical composites [7], are not evolved after oxidation in the present study; the  and SiC-depleted layers  for oxidized ZS  formed in the  are  detection of Si with O is much lower than that of Zr. The ZrO2 layer formed on the ZS10 specimen is thinner than that on the monolithic  ZrB2. However, the average particle size in the ZrO2 layer for the ZS10 specimen (7 µm) is larger than that for the monolithic ZrB2 specimen (6 µm). The SiC-depleted layer has relatively lower porosity and larger  ZrB2 grain size in ZS10 than in the other ZS specimens. These results indicate that the evolution of SiO by the oxidation of SiC and the  coarsening of ZrB2 particles occur For the ZS20 and ZS30 specimens, ZrO2 and SiC-depleted layers are formed, but SiO2 and ZrO2+SiO2 layers are not, as with ZS10.  simultaneously during oxidation.  8085  Fig. 8. The changes in thickness of each layer as a function of the volume fraction of SiC  particles.  In the present study, the thickness of the oxidized region, hoxi, was defined as  h  oxi  =  h  ZrO2  +  h  SiC− dep  (4)  where hZrO2 and hSiC-dep indicate the thicknesses of the dense ZrO2 and SiC-depleted layers, respectively. Fig. 8 shows the changes in thickness  of each layer as a function of the volume fraction of SiC particles. The  oxidized layer thickness (hoxi) is increased with increases in the volume fraction of SiC, while the thickness of the ZrO2 layer (hZrO2) remains nearly constant among the oxidized specimens of monolithic ZrB2 and ZS composites in the present study. The morphologies of the ZrO2 layers are also similar for the ZS composites; however, the thicknesses  of the ZrO2 and SiC-depleted layers on ZS20 are the lowest among the specimens oxidized in the present study. Fig. 9 shows the relationship  between the porosity of  the SiC-depleted layer and the SiC content  in  the ZS composites. The results obviously indicate that the SiC-depleted  layer porosity is increased with increases in the SiC content.  4. Discussion  4.1. Composition and morphologies after heat exposure up to 2000 °C  Under  the present  testing conditions,  the SiC particle-dispersed  ZrB2 matrix composites formed porous ZrO2 layers on the surfaces. It is well known that ZrO2 is formed by the oxidation of ZrB2 [27]:  \\x0c', 'ZrB2(s) + 5/2 O2(g) → ZrO2(s) + B2O3(l)  (5)  B2O3(l) → B2O3(g)  (6)  B2O3(l) should evaporate at the present testing temperature of over 2000 °C because of the high vapor pressure of the liquid [27] In  addition, a 16% volume expansion occurs because of the difference in density between ZrB2 (6 g/cm3) and ZrO2 (5.68 g/cm3)[28]. Generally, ZS composites form several-layered structures of SiC depleted, SiO2-filled ZrO2, and porous ZrO2 layers under oxidation at temperatures reaching 1800 °C in air [28]. Two possible reactions by  the oxidation of SiC may occur as follows [29]:  SiC(s) + 3/2 O2(g) → SiO2(l) + CO(g)  (7)  SiC(s) + O2(g) → SiO(g) + CO(g)  (8)  SiO(g) + 1/2O2(g) →SiO2(l)  (9)  Reactions (7) and (8) are called “passive” and “active” oxidations of  SiC, respectively, and their occurrences depend on the temperature and  oxygen partial pressure. In addition, the formation and decomposition  of SiO2 occurs by reaction (9). For ZS composites, a high oxygen partial  pressure exceeding ~105 Pa is  required to form SiO2 layer during oxidation (by reactions (7) or (9)); this is higher than the  in an oxidized  oxygen partial pressure in ambient air [29]. However, gaseous SiO is  formed, and SiO2 forms not at smoke, because SiO reacts with oxygen to form SiO2 above the surface because the temperature is lower there than it is at the surface [29]. In  the surface but over  the surface as a  the present study, the oxidation test was conducted in ambient air with an estimated oxygen partial pressure of approximately 2.0×104 Pa. The  oxygen partial pressure decreases locally with the active oxidation of  SiC [30]. SiO2 is formed when the temperature decreases to less than 1825 °C [31]. Thus, for the ZS composites, the formation of SiO2 also occurs over the surface, and neither SiO2 scale nor ZrO2 + SiO2 scale is formed in the oxidized region. At temperatures above 2000 °C, gaseous  SiO is formed by the active oxidation of SiC and re-oxidized over the  specimen to form SiO2 as smoke. In static conditions, the accumulation of SiO2 on the specimen surface, formed by the oxidation of SiO, has been reported [32]. Thus, it is  suggested that SiO2 specimen during the  formed by  the  oxidation  of  SiO covers  the  oxidation  test.  Furthermore, Fig. 5 shows  that neither an SiO2 nor an SiO2+ZrO2 is formed during oxidation. This result affects the temperature  layer  measurement by the  IR thermometer. As noted in Section 3.1,  the  emissivity values of the ZS composites vary during oxidation. Although  some researchers have reported the emissivity values of ZS composites  during oxidation below 2000 °C, SiO2 or SiO2+ZrO2 layers covered the surfaces of these ZS composites [23,24]. In the present study, only  ZrO2 surface. The ZrO2 (wavelength: 0.65 µm) is 0.62-0.75 [33], lower than the reported emissivity of oxidizing ZS composites (0.7-0.95). Thus, the maximum  is  formed  on  the  emissivity  of  for  IR  surface  temperature  is  estimated  as ~3000 °C from Eq.  (3).  It  is  suggested that the temperature measurement of the oxidation test for  ZS  composites  by  the  IR thermometer  depends  strongly  on  the  compounds covering the surfaces of  the ZS composites during oxida tion. The  emissivity  of ZrO2 oxidation above 2000 °C, because only porous ZrO2 layers are formed on the surfaces by oxidation in this temperature regime.  in particular must  be  considered for  Fig. 10 shows schematics of  the oxidation mechanisms of mono lithic ZrB2 and ZS composites, based on the above-mentioned experimental results. For the monolithic ZrB2 specimen, a ZrO2 layer with a porous structure is formed after the oxidation test (see Figs. 3 and 4).  However,  this layer is dense because the ZrO2 particles form necking structures evolved by sintering, since the sintering of monolithic ZrO2  0  30  P  o  r  o  s  i  t  y  f  o  r  S  i  C   p e d  l  e  t  d e  l  a  y  e  r  (  %  )  20  10  5  15  25  0  10 20 30 Content of SiC in ZS composites (vol%)  40  Fig. 9. Thicknesses of porous ZrO2 layers and SiC-depleted layers as functions of SiC  contents.  Fig. 10. Schematics of oxidized monolithic ZrB2 and ZS composites.  R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  8086  \\x0c', 'begins above 1200-1400 °C [34,35] and volume expansion occurs. On  the other hand, gaseous B2O3 supplied by the oxidation of ZrB2 during the oxidation test (Eq. (1)). Paths for the escape of gaseous B2O3 are formed continuously in the ZrO2 layer during oxidation. Thus, the sintering of the ZrO2 layer does not continue because the distance between ZrO2 particles cannot shrink to close the paths for B2O3. For the ZS composites, porous ZrO2 layers and SiC-depleted layers are formed after oxidation testing. The microstructure of the porous ZrO2 layers formed on the ZS composites differs from that on the monolithic  is  ZrB2. Skeletons of ZrO2 are connected continuously for the layers on ZrB2, whereas the ZrO2 particles are separated by cracks and voids in the layers on the ZS composites. This difference is attributed to the  amounts of gaseous products evolved during the oxidation tests. B2O3 is the only gaseous product of the oxidation of monolithic ZrB2. for the ZS composites, SiO and CO are also evolved during  However,  oxidation,  in addition to B2O3. This implies that the gaseous products evolved during oxidation prevent the sintering of the ZrO2 layer. In the SiC-depleted layer becomes porous as shown in Fig. 4  addition,  because  the  active  oxidation  of  SiC (Reaction  (8))  only  produces  gaseous products.  In the present study,  it seems that 20 vol% of SiC  particles effectively prevents  the formation of a porous SiC-depleted  layer.  4.2. Thickness of SiC-depleted layer  In the present  study,  the SiC-depleted regions were  formed as  oxidized areas in the ZS composites. The formation was caused by the  preferential oxidation of SiC compared to ZrB2, as with other results from conventional oxidation tests. A low oxygen partial pressure,  sufficient  to  evolve  gaseous SiO by  the  active  oxidation of SiC,  is  induced by the low oxygen diffusivity of the SiO2 and ZrO2+SiO2 layers. We summarize the thicknesses of SiC-depleted layers reported by some  researchers in Fig. 11 [2,13,14,36,37]. This figure shows the relation ship between the thickness of the SiC-depleted layer and the oxidation  time; the behavior of the graph differs completely based on whether the  oxidation temperature reaches 2000 °C. For oxidation above 2000 °C,  the thickness of  the SiC-depleted layer is increased with increases in  the oxidation time. For oxidation below 2000 °C, the layer thickness is  independent of  the oxidation time. This  suggests  that  the oxidation  mechanism of ZS composites  above 2000 °C is different  from that  below 2000 °C. ZS composites form ZrO2+SiO2 scale during oxidation below 2000 °C; the SiO2 prevents oxygen diffusion toward the unoxidized regions of the composites. For oxidation tests below 2000 °C, the  active oxidation of SiC occurs because the area below the ZrO2+SiO2 scale develops a low oxygen partial pressure (4.0×10−14−1.8×10−11 at  1500 °C [7]) and oxidation proceeds as a diffusion-limited reaction  because of  the low oxygen permeability of SiO2. On the contrary, as above, ZS composites do not form SiO2 scale during oxidation above 2000 °C. SiC reacts with oxygen to form gaseous SiO  mentioned  above 1825 °C because the oxygen partial pressure of atmospheric air (2.0×104 Pa)  is  insufficient  to  form SiO2 Furthermore, Fig. 11 also clearly indicates that the thickness of the SiC by  the  oxidation of  SiC.  depleted layer formed by oxidation above 2000 °C for ~10 s is almost the same as that formed by oxidation below 2000 °C for ~5-60 min and is ~5-10% of formed by oxidation above 2000 °C for ~5-  that  15 min. This  suggests  that  the SiC-depleted layer  is  formed at  the  beginning  of  oxidation, within ~10 s  of  the  temperature  reaching  2000 °C, and that  the formation of  the SiC-depleted layer  is unique  to the rapid oxidation of ZS by fast heating.  In addition,  the heating  rate of the oxidation test in this study is more than ~300 °C/s, much higher than that in furnace testing (0.1-0.5 °C/s) [1,5], arc-jet testing  (~50 °C/s)  [10],  and  torch  testing  (~50 °C/s)  [14].  Furthermore,  resistance heating can be conducted in still air, whereas  torch tests  and arc-jet tests are affected by the dynamic pressure. It is implied that  the  SiC-depleted layers  are  formed because  of  the  extremely high  temperature (above 2000 °C). Thus,  the effect of  temperature on the  oxidation of ZS composites under high heating rates can be extracted  from the present study.  Based on Figs. 5 and 8, the thickness of the SiC-depleted region for  ZS20 was the lowest among the materials oxidized in this study. As  mentioned above, ZrB2 begins sintering of ZrB2 particles in the SiC-depleted layer occurs neously with the sintering  to sinter above 2000 °C [20] and the  simulta of  the unoxidized region. However,  the  porosity of the SiC-depleted region increased with increases in the SiC  content  because  the  preferential  oxidation  of  SiC  (reaction  (5))  produced gaseous products, and voids formed where SiC had existed  prior to oxidation. Thus, the distance between ZrB2 particles in the SiCdepleted layers was increased and the sintering of ZrB2 was further suppressed with the increasing of SiC content in the ZS composites. In  the present study, ZS20 formed a thinner oxidized layer than ZrB2, and the thickness of the SiC-depleted layer in ZS20 was the lowest among  the ZS composites. These results indicated that  the increasing of  the  SiC content  in ZS composites caused increases in the porosity for the  SiC-depleted layer. Meanwhile, the decreasing of the SiC content in the  ZS composites caused increases in the thickness of  the SiC-depleted  layers. Then, below the SiC-depleted layers, the oxygen partial pressure  was decreased and the oxidation of SiC did not occur even under active  oxidation conditions. Therefore, the unoxidized regions remained after  oxidation  testing.  The  decreases  in  the  SiC  contents  in  the  ZS  composites  caused decreases  in the  consumption of oxygen by  the  oxidation of SiC;  therefore, oxygen penetration into the unoxidized  region was increased at lower SiC levels. Thus, in the present study, the  addition of 20 vol% SiC to ZrB2 oxidation of ZS composites.  is optimal  for preventing the rapid  5. Conclusion  In the present study,  the oxidation resistance of monolithic ZrB2 composites with different SiC contents was successfully  and  ZS  evaluated  above  2000 °C at  high  heating  rates  using  a  resistance  heating system. The following results and conclusions were obtained:  1. A resistance heating system was  successfully applied in oxidation  testing of ZrB2 and ZS composites with different SiC contents above 2000 °C. The ZS composites formed oxidized layers comprising  columnar-shaped ZrO2 layers and SiC-depleted layers. The morphologies of the oxidized layers strongly depended on the SiC contents.  2. The  formation mechanism of  the  SiC-depleted  layers  differed  dramatically based on whether the oxidation temperature exceeded  2000 °C. At  temperatures of or above 2000 °C, SiC-depleted layers  were formed because the oxygen partial pressure of  the air atmo sphere was not enough to form SiO2 by the oxidation of SiC.  Fig. 11. Plots of  literature and experimental data of SiC-depleted layer thickness as a  function of oxidation time [2,13,14,36,37].  R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  8087  \\x0c', 'R. Inoue et al.  Ceramics International 43 (2017) 8081-8088  3. The porosity of  the SiC-depleted layer was  increased as  the SiC  content  in the ZS composites was increased because of  the greater  amount  of  SiC  particles  consumed  in  the  SiC-depleted  layer.  However,  the thickness of  the SiC-depleted layer was increased as  the SiC content  in the ZS composites was decreased, because the  decreasing of  the SiC content  increased the amount of oxygen that  penetrated the unoxidized region. Although the porosities of the SiC depleted layers were intermediate in the ZS composites oxidized in  the present study, ZS20 showed the best oxidation resistance above  2000 °C  at  high  heating  rates  because  ZS20  had  the  thinnest  oxidized region and SiC-depleted layer. 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Hilmas, Pressureless sintering of zirconium diboride, J. Am. Ceram. Soc. 89 (2006) 450-456. [26] M. Ishigame, T. Sakurai, Temperature dependence of the Raman spectra of ZrO2, J. Am. Ceram. Soc. 60 (1977) 367-369. T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, A model for the oxidation of ZrB2, HfB2 and TiB2, Acta Mater. 55 (2007) 5999-6010. [28] W.G. Fahrenholtz, The ZrB2 volatility diagram, J. Am. Ceram. Soc. 88 (2005) 3509-3512. [29] A. Rezaie, W.G. Farenholtz, G.E. Hilmas, Evolution of structure during the  [23]  [27]  [24]  oxidation of zirconium diboride-silicon carbide in air up to 1500 °C, J. Eur. Ceram. Soc. 27 (2007) 2495-2501. [30] N.S. Jacobson, D.L. Myers, Active oxidation of SiC, Oxid. Met. 75 (2010) 1-25. [31] Y. Kubota, H. Hatta, T. Yoshinaka, Y. Kogo, T. Goto, T. Rong, Use of volume  [32]  element methods to understand experimental differences in active/passive transitions and active oxidation rates for SiC, J. Am. Ceram. Soc. 96 (2013) 1317-1323. J.W. Hinze, H.C. Graham, The active oxidation of Si and SiC in the viscous gas‐flow regime, J. Electrochem. Soc. 123 (1976) 1066-1073. [33] D. Cubicciotti, The melting point—composition diagram of the zirconium—oxygen system, J. Am. Chem. Soc. 73 (1951) 2032-2035. [34] K.T. Kim, H.G. Kim, H.M. Jang, Densification behavior and grain growth of  zirconia powder compact under high temperature, Int. J. Eng. Sci. 36 (1998) 1295-1312. F. Wakai, T. Nagono, The role of  interface-controlled diffusion creep on super [35]  plasticity of yttria-stabilized tetragonal ZrO2 polycrystals, J. Mater. Sci. Lett. 7 (1988) 607-609. [36] X. Li, X. Zhang, J. Han, C. Hong, W. Han, A technique for ultrahigh temperature oxidation studies of ZrB2-SiC, Mater. Lett. 62 (2008) 2848-2850. F. Monteverde, R. Savino, Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions, J. Eur. Ceram. Soc. 27 (2007) 4797-4805.  [37]  8088  \\x0c']"
},{
  "_id": 156,
  "PDF": "Oxidation behaviour at 1100°C in air of 25 wt-_ Cr-containing cobalt-based alloys containing high quantities of hafnium carbides.pdf",
  "Text": "['Canadian Metallurgical Quarterly  The Canadian Journal of Metallurgy and Materials Science  ISSN: 0008-4433 (Print) 1879-1395 (Online) Journal homepage: https://www.tandfonline.com/loi/ycmq20  Oxidation behaviour at 1100°C in air of 25\\u2005wt-% Cr-containing cobalt-based alloys containing high quantities of hafnium carbides  P. Berthod & E. Conrath  To cite this article: P. Berthod & E. Conrath (2016) Oxidation behaviour at 1100°C in air of 25\\u2005wt% Cr-containing cobalt-based alloys containing high quantities of hafnium carbides, Canadian Metallurgical Quarterly, 55:4, 409-419, DOI: 10.1080/00084433.2016.1206290  To link to this article:  https://doi.org/10.1080/00084433.2016.1206290  Published online: 19 Jul 2016.  Submit your article to this journal   Article views: 63  View related articles   View Crossmark data  Citing articles: 2 View citing articles   Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=ycmq20  \\x0c', 'Oxidation behaviour at 1100°C in air of 25 wt-% Cr-containing cobalt-based alloys containing high quantities of hafnium carbides  P. Berthod* and E. Conrath  Hafnium, usually added to improve the high temperature oxidation resistance of alloys, allowed  obtaining HfC carbides which are very efficient  for  the creep-resistance. For  that Hf must be  added in particularly high quantities which may possibly  influence the oxidation behaviour.  Three HfC-strengthened cast cobalt alloys were studied all along thermogravimetry  tests at  1100°C. They were compared to similar but Hf-free ternary alloys. The mass plotted according to {m × dm/dt = f (-m)}  to specify all kinetic oxidation constants, and versus  variation were  temperature to study the oxidation beginning during heating and the scale spallation during  cooling.  The  presence  of many HfC carbides  obviously  influences  the  high  temperature  oxidation: mass gain occurring sooner during heating,  faster  isothermal mass gains but better  behaviour  in oxide scale spallation during cooling. This deterioration of oxidation behaviour  must be corrected to hope benefiting from the high creep-resistance brought by this new type  of strengthening.  Le hafnium, habituellement ajouté pour améliorer la résistance à l’oxydation à haute température  des alliages, a permis d’obtenir des carbures HfC qui sont  très efficaces pour  la résistance au  fluage. Pour cela, on doit ajouter  le Hf  en quantités particulièrement  élevées, ce qui peut  influencer  le comportement à l’oxydation. On a étudié trois alliages de cobalt coulés renforcés  par des carbures HfC, au moyen d’essais de thermogravimétrie à 1100°C. On les a comparés à  des alliages ternaires similaires, mais sans Hf. On a tracé une courbe de la variation de la  masse  d’après  {m × dm/dt = f(-m)}  pour  spécifier  toutes  les  constantes  de  la  cinétique  d’oxydation, en fonction de la température, afin d’étudier le début d’oxydation lors du chauffage  et  l’écaillage de  l’oxyde  lors du  refroidissement.  La présence de  nombreux  carbures HfC  influence évidemment  l’oxydation à haute température : gain de masse se produisant plus tôt  lors  du  chauffage,  gains  de masse  isothermes  plus  rapides mais meilleure  résistance  à  l’écaillage  d’oxyde  lors  du  refroidissement. On  se  doit  de  corriger  cette  détérioration  du  comportement à l’oxydation si  l’on espère bénéficier de la haute résistance au fluage, amenée  par ce nouveau type de renforcement.  Department Chemistry and Physics of Solids and Surfaces, Team ‘Surface and Interfaces, Chemical Reactivity of Materials’, Faculty of Sciences and Technologies, University of Lorraine, Institut Jean Lamour (UMR CNRS 7198), Boulevard des Aiguillettes, B.P. 70239, 54506 Vandoeuvre-lèsNancy, France  *Corresponding author, email patrice.berthod@univ-lorraine.fr  © 2016 Canadian Institute of Mining, Metallurgy and Petroleum Published by Taylor & Francis on behalf of the Institute Received 30 December 2015; accepted 20 June 2016 DOI 10.1080/00084433.2016.1206290  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  409  \\x0c', 'Be r thod and Conra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  Keywords: Cobalt-based alloys, Hafnium carbides,  Isothermal oxidation, Oxidation at heating, Oxide spallation at cooling, Oxide characterisation  Introduction  Many high temperature applications require geometrically complicated metallic pieces and the classical foundry way is often essential to allow complex-shaping1 as well as coarse-grained microstructures to keep high creep-resistance. The obtained polycrystalline alloys present multiple randomly oriented grain boundaries which are the mechanically weakest zones in the microstructures. When the alloys must be used under stresses at temperatures higher than the equi-cohesive temperature, their grain boundaries must be reinforced by hard particles. These ones can be the eutectic carbides which appeared at the end of solidification in the interdendritic spaces. This is the case of the acicular chromium carbides or script-like tantalum carbides which are often exploited in carbon-containing cast cobalt-based superalloys.2 The carbides mechanically strengthening grain boundaries must be very stable at the high temperatures of work. Chromium carbides do not meet this requirement since their volume fractions rapidly decrease at high temperature. More stable, the tantalum carbides are nevertheless sometimes not heatresisting enough since they get fragmented too rapidly3 with as result significant loss in high temperature mechanical resistance.4-6 Other MC carbides (i.e. monocarbide of a ‘M’ metal, ‘M’ being Ta for example) are particularly resistant to morphological changes at high temperature: the HfC carbides recently obtained in solidified cobalt-based alloys7 with a script-like shape and forming an interdendritic eutectic compound with matrix. For several decades hafnium was mainly considered as an element beneficial for the high temperature oxidation behaviour of superalloys. Its actions in this field are slowing the oxide growth or improving the protective oxides adherence.8 According to recent articles, the used Hf contents were generally close to 1 wt-% (typically between 0.25 and 1.5 wt-%) in commercial alloys,9 iron-based bulk alloys rich in Cr and Al,10 aluminides11 and TBCsupporting bond coats.12 Hafnium was recently used in alloys for use at high temperature in order to develop a particularly efficient strengthening interdendritic network of HfC carbides. Their mechanical properties at temperatures up to 1200°C were remarkable for cast polycrystalline alloys simply reinforced by carbides.13 However, alloys containing so high Hf contents, an element which is particularly easy to oxidise at high temperature, may present particular behaviours in oxidation at high temperature, which may be interesting to explore. This is what was undertaken in the present study for 1100°C, temperature at which the best actual superalloys - the γ/γ′ Ni-based single-crystals lose their principal strengthening particles, with three cast Co-based alloys containing various volume fractions in HfC carbides. In parallel two cobalt-based ternary alloys with the same contents in chromium and carbon but without hafnium were also studied in order to specify the base behaviour  -  410  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  and then, by comparison, carbides.  the influence of  the hafnium  Experimental methods  Alloys’ elaboration  .  .  .  .  (atomic  (C atomic  For each alloy a total charge of about 40 g of pure elements (>99.9 wt-%, Alfa Aesar) was prepared, according to the following compositions (in wt-%): ‘CoHf1’ alloy: Co(bal.)-25Cr-0.25C-3.7Hf equality between C and Hf) ‘CoHf2a’ alloy: Co(bal.)-25Cr-0.5C-3.7Hf content equal to twice the Hf one) ‘CoHf2b’ alloy: Co(bal.)-25Cr-0.5C-7.4Hf (atomic equality between C and Hf again, double HfC population expected by comparison with the CoHf1 alloy) ‘Co1’ alloy: Co-(bal.)-25Cr-0.25C (to which CoHfC1 alloy will be compared) . ‘Co-2’: Co(bal.)-25Cr-0.50C (to which the CoHf2a and CoHf2b alloys will be compared). The elaboration of these alloys, whose compositions are summarised in Table 1, was carried out using a High Frequency induction furnace (CELES, France; operating parameters: voltage 4 kV and frequency 110 kHz). Melting, the 3 min-dwell and solidification of the alloys took place in a water-cooled copper crucible present in the fusion chamber of the furnace, in a surrounding inert gaseous environment (300 mbars of pure argon). Compact ingots were obtained, and after complete cooling, they were cut to obtain the parts needed for the characterisation of the as-cast microstructures of the alloys as well as the one destined to the high temperature oxidation tests.  the  Thermogravimetry tests and kinetic exploitations  The parts destined to the high temperature oxidation tests were cut to obtain parallelepipeds of about 8 × 8 × 3 mm3. These ones were ground, first by using 240-grit SiC papers to smooth edges and corners, and second all around with 1200-grit SiC papers. A sample of each of the five alloys was subjected to high temperature oxidation in a thermo-balance. The used apparatus was a SETARAM TGA92 one. Tests were done in a constant flow of synthetic air (80% N2- 20% O2). In each case, after a heating at +20°C min−1,  Table 1  Chemical compositions of the five studied alloys (all contents in wt-%)  Elements  Co1 CoHf1 Co2 CoHf2a CoHf2b  Cr  25 25 25 25 25  C (expected)  0.25 0.25 0.50 0.50 0.50  Hf  / 3.7 / 3.7 7.4  \\x0c', 'the samples were maintained at 1100°C during 50 h, then cooled at −5°C min−1, rate expected low enough to limit the spallation of the oxide scale developed all over the sample. The mass gain files (record step: about 40 s) were thereafter exploited with a spreadsheet program. It was first verified, by plotting time versus the mass gain recorded during the 1100°C-stage, that the isothermal oxidation kinetic was globally parabolic type and then that it seemingly follows the Wagner’s law. This parabolic regime may be preceded by an initial short linear transient oxidation, the linear constant of which, Kl, was determined. The isothermal part of the mass gain files was first classically treated to determine the parabolic constant Kp, by plotting mass gain versus the square root of time and by noting the average slope of the obtained more or less straight curve. Since the studied alloys possibly present a chromiaforming behaviour thanks to the 25 wt-%Cr that they contain, a minimisation of the oxidation rate when assessed from mass gain results may occur because of the possible mass loss by chromia volatilisation. This phenomenon, re-oxidation of Cr2O3 in volatile CrO3 is expected in air when temperature is higher than 1000° C. Even at 1100°C chromia volatilisation must be taken into account, by determining the chromia volatilisation constant Kv together with Kp. For this purpose one used a method recently developed.14 This one con{−m} sists in plotting, versus (opposite of the mass gain per surface unit area), the product (the mass gain per surface unit area multiplied by the same quantity but derived by regard to time). After a first transient part, the plotted versus {−m} becomes a straight line obeying the {m × dm/dt equation. The parabolic constant Kp and the chromia volatilisation constant Kv can be thus easily deduced as, respectively, the ordinate at the origin and the slope of the straight line. Before that, the mass gain during the heating (performed in presence of synthetic air) was specified. This first mass gain needs to be known and added to the eventual linear mass gain to allow an accurate determination of Kp since the oxide thickness already existing at the beginning of the isothermal stage influences the parabolic regime. Since the used thermo-balance was not a symmetric one, the mass gain at heating and the one at cooling were corrected from the air buoyancy variations as done in an earlier work.15 For this the mass gains were plotted versus temperature instead of time, and corrected from the apparent mass variations due to the buoyancy variations of the heated air. The corrected curves of mass gain versus temperature during the heating allow specifying first the temperature at which the mass gain really due to oxidation becomes to be detectable by the used thermo-balance. This criterion is of course not absolute since it closely depends on the used thermo-balance. However, the comparison of these ‘oxidation start’ temperatures between the different alloys oxidised in the same thermo-balance and for the same oxidising atmosphere and heating rate, may be very interesting.  = Kp - m × Kv}  (or Kp - Kv)  {m × dm/dt}  {m × dm/dt}  Be r thod and Con ra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  The cooling parts of these curves of mass gain, corrected from the air buoyancy variations and plotted versus temperature, can give indications about oxide spallation.16 The deduced parameters were the temperatures at which oxide spallation started and the final mass variation during the whole cycle.  Metallography characterisation  The parts destined to the observations of the as-cast microstructures were embedded in a cold mixture composed of a liquid resin (DBF from ESCIL, 82% in mass) and a hardener (HY956 from ESCIL, 18%). They were thereafter ground, using first SiC papers with grades from 120 or 240 up to 1200 under water, and second, after intermediate ultra-sonic cleaning, polished using a textile disk enriched with diamond particles (ESCIL) to obtain a mirror-like surface state. After the oxidation tests the samples were first subjected to X-ray diffraction analysis (Philips X’Pert Pro MPD diffractometer, copper anode, wavelength λ = 1.5406 Å) to identify the oxides formed on surface. Second, they were coated with a fine gold layer to give electric conductivity to their surface (despite the presence of external oxide scales). This was done to allow thereafter an electrolytic deposition of a thick nickel shell destined to protect the oxide scales from loss during cutting. The gold-covered oxidised samples were immersed in a warm (50°C) Watts’ bath (aqueous solution of Ni2+ ions prepared from mainly NiSO4 and NiCl2 salts) to allow the deposition of metallic nickel by reduction of the Ni2+ cations under a current density of about 1.6 A dm−2. After cutting the oxidised samples in two parts, cross-sections were prepared by following the same metallographic preparation details (embedding- grinding-polishing) as for the samples for the as-cast microstructure examinations. The global chemical composition of each as-cast alloy was controlled using the energy dispersive spectrometry (EDS) device attached to a scanning electron microscope (SEM) (model: JSM6010LA from JEOL), under an acceleration voltage of 20 kV The initial microstructures of the as-cast alloys were examined using the SEM, mainly in the back scattered electrons mode (BSE). The oxide scales and the subsurface microstructures (in cross-section) were also examined by SEM/BSE and analysed by spot SEM/EDS measurements. Additionally, successive spot EDS analyses were performed in the alloy at an increasing depth from the oxide/alloy interface, perpendicularly to later one. This aimed to specify the concentration profiles in Cr and Hf in the subsurface, in order to better describe the deterioration of the external part of the alloys due to oxidation.  Results and discussion  As-cast microstructures of the alloys  The microstructures of the three Hf-rich alloys in their ascast conditions are illustrated in Fig. 1 (SEM images taken in BSE mode). They present an interdendritic  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  411  \\x0c', 'Be r thod and Conra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  1  As-cast microstructures of the three Hf-rich alloys (SEM/BSE examinations)  network of script-like HfC carbides. These ones are in the CoHf2a alloy as numerous as in the CoHf1 one, while they are much more present in the CoHf2b alloy. In addition to the interdendritic script-like HfC carbides one can note the presence of other HfC carbides, geometrically different. These ones, more compact, have precipitated at the beginning solidification: they are pro-eutectic HfC carbides while the script-like HfC ones have precipitated as eutectic with matrix. The volume fraction of HfC carbides of both types is obviously closely dependent on the Hf content in the alloy, and not on the carbon content. Dark particles are additionally present in the CoHf2a alloy. They are chromium carbides. The two ternary alloys are composed of matrix and of chromium carbides. These carbides are however rather rare in the Co1 alloy. In contrast, they are more present in the Co2 one (higher carbon content), in the interdendritic spaces. For a given carbon content the presence of hafnium obviously leads to much more carbides than in the ternary alloys. This shows that Hf is a much stronger carbides-forming element. Some pinpoint EDS measurements performed in the matrix showed that hafnium is totally absent in the matrix of the Hf-rich alloys. This indicates that all the hafnium atoms have formed HfC.  Isothermal oxidation  . Kl  Hf-rich one was a long time almost linear before becoming more parabolic, and the Hf-free one was linear during a short time, thereafter parabolic for about 30 h until mass gain accelerated again. The three 0.50C-containing alloys (Fig. 2 bottom) led to much more parabolic curves than the two previous ones. The oxidation rates seem to be higher in presence of hafnium, and the higher the Hf content the faster the mass gain rate. This is confirmed by the results of kinetic exploitation of the mass gain files (Table 2): constants: rather scattered results (8-40 × 10−8 g cm−2 s−1), no evident dependence on the carbon and hafnium contents . Kp constants (classically determined): parabolic kinetics faster in presence of Hf (from about 80 to about 270 × 15 × 10−12 g2 cm−4 s−1 against about for the two Hf-free alloys); higher Kp values noticed in presence of only HfC (CoHf1 and CoHf2b), lower values in presence of chromium carbides (Co1, Co2 and even CoHf2a which contains a little chromium carbides in addition to HfC) . Kv constants: significantly higher for the Hf-rich alloys (from about 270 to about 400 × 10−10 g cm−2 s−1, against 40-80 × 10−10 g cm−2 s−1 for Co1 and Co2); this does not change the Kp hierarchy between the Hf-free alloys (from about 20 to about 40 × 10−12 g2 cm−4 s−1) and the Hf-rich ones (from about 240 to about 470 × 10−12 g2 cm−4 s−1).  10−12 g2 cm−4 s−1,  The mass gain curves plotted versus time are all displayed in Fig. 2. The two 0.25C-containing alloys (Fig. 2 top) led to no really parabolic curves. Indeed, the mass gain of the  The values of kinetic constants Kp and Kv, particularly high for the Hf-rich alloys, were however accurately determined and validated by comparison between  412  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  \\x0c', 'Be r thod and Con ra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  3  Example of mass gain values plot according to the {m × dm/dt = f (−m) = Kp - m × Kv} to specify the isothermal kinetic constants Kp and Kv (here: the CoHf2a alloy)  2  Mass gain curves recorded for the three Hf-rich alloys with  comparison with the corresponding Hf-free ternary alloys  (top: 0.25C-containing alloys, bottom: 0.50C-containing  alloys)  experimental curve and mathematical ones plotted from the determined values of Kp and Kv. The {m × dm/dt = f (−m) = Kp - m × Kv} method well led to really straight lines the equation of which allowed determining the two constants (e.g. Fig. 3), which were validated by a very faithful mathematical curve (green dotted line in Fig. 4, much closer to the thick grey experimental curve than the thin red mathematical curve plotted with the Kp value classically determined).  4  Comparison, with  the  experimental mass  gain  curve  (‘exp’), of the mathematical curve drawn using the value of Kp only (issued from classical determination of parabolic constant, ‘math Kp classic’) and of the mathematical curve drawn using the values of Kp and Kv (issued from the treatment according to {m × dm/dt = f (−m) = Kp - m × Kv}, ‘math Kp & Kv’); (here: the CoHf2a alloy)  Oxidation at heating and scale spallation at cooling  The mass gains during heating are plotted versus temperature in Fig. 5 (top) for the 0.25C-containing alloys and in Fig. 5 (bottom) for the 0.50C-containing ones. One can see that the mass gain by oxidation became detectable by the thermo-balance sooner for a lower carbon content (i.e. at lower temperature) than for a higher one, and for a Hf-rich alloy than for a Hf-free one. These observations, confirmed in Table 3 by the values  Table 2  Values of the kinetic constants determined from the mass gain files  Constants Alloys  Co1 CoHf1 Co2 CoHf2a CoHf2b  Kl  (×10−8 g cm−2 s−1) Slope at stage start  7.73 19.5 39.2 15.7 29.6  Classic  12.4 253 15.4 76.0 274  Kp (×10−12 g2 cm−4 s−1)  m × dm/dt = f (−m)  Kv (×10−10 g cm−2 s−1)  m × dm/dt = f (−m)  19.4 433 41.4 242 470  40.1 267 80.8 395 268  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  413  \\x0c', 'Be r thod and Conra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  5  Start of mass gain by oxidation during the heating (after  correction of the thermogravimetry files from the air buoy ancy variations); top: 0.25C-containing alloys alloys (with  inserted enlargement of the oxidation start part), bottom:  0.50C-containing alloys  of oxidation start temperature, let think first that a more dense interdendritic carbide network tends favouring an early oxidation during heating by exposing more the highly oxidable elements (Cr, Hf) to hot air, and second a more oxidable carbide-former element leads to the  6  Start of oxide scale spallation during the cooling (after  correction of the thermogravimetry files from the air buoy ancy variations);  top: 0.25C-containing alloys, bottom:  0.50C-containing alloys  same effect by a better reactivity with oxygen. In contrast, the total mass gain achieved during the whole heating does not follow a clear law versus the carbon and hafnium contents (Table 3). Plotting mass gain or variation versus temperature allows evidencing the moment at which the oxide scales, previously formed during heating and mainly during the isothermal stage, started to be lost (Fig. 6). These curves, as well as the values resulting of their exploitation and  Table 3  Oxidation behaviour during heating: temperature of oxidation start (according to the accuracy of the used thermo-balance) and total mass gain achieved by oxidation before the isothermal stage beginning  Table 4  Oxidation behaviour during cooling: temperature of scale spallation start and final mass variation after the whole thermal cycle  Alloys  Co1 CoHf1 Co2 CoHf2a CoHf2b  Temperature of oxidation start (°C)  Mass gain during heating (mg cm−2)  844 819 932 847 859  0.056 0.089 0.098 0.069 0.136  Alloys  Co1 CoHf1 Co2 CoHf2a CoHf2b  Temperature of spallation start (°C)  Final mass variation (g cm−2)  578 385 460 416 340  −0.004 +0.003 −0.003 0.000 +0.006  414  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  \\x0c', 'Be r thod and Con ra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  consequently of high temperature mechanical properties upholder, hafnium remained efficient as scale adherence favouring element, despite its particularly high contents here.  Pre-cutting oxide scale characterisation  The surfaces of the oxidised samples are illustrated in Fig. 7 by macrographs taken using a simple office scanner. They confirm that the scales formed on all alloys tended to spall off during cooling but also that final spallation state was more severe for the Hf-free alloys than for the Hf-rich ones. XRD measurements were carried out on the oxidised surfaces of the samples before they were cut. Results (Table 5) showed that chromia (Cr2O3) was present at the surface of all alloys, but it was never the single oxide present. The CoCr2O4 spinel oxide, and even the simple CoO cobalt oxide were often present too. The HfO2 oxide was also detected on the oxidised surfaces of the three Hf-rich alloys.  Cross-sectional characterisation of oxide scales and subsurface deterioration  The thickness of locations of the  the scales as well as the amounts and different oxides revealed by XRD  7  Scanned views of one of  the faces of  the samples after  oxidation test, for illustrating their oxidised states  which are displayed in Table 4, show that the spallation of the external oxide scales occurred in all cases, but it started after a cooling down to a temperature which was significantly lower for the alloys containing hafnium than for the Hf-free ternary ones. As a result the total mass change is positive for the Hf-containing alloys and negative for the ternary ones. Thus, besides its principal role of carbides stabiliser at high temperature and  8  Illustration of the surface states of the two 0.25C-containing alloys after oxidation test (cross-sectional SEM/BSE examin ations, SEM/EDS oxide determinations); top: the Hf-free Co1 alloy, bottom: the Hf-rich CoHf1 alloy  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  415  \\x0c', 'Be r thod and Conra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  Table 5  Natures of the various oxides present in the external scales (XRD results)  Alloys  Co1 CoHf1 Co2 CoHf2a CoHf2b  Composition of the oxide scales  CoO-CoCr2O4-Cr2O3 CoO-CoCr2O4-Cr2O3 CoO-Cr2O3 CoCr2O4-Cr2O3-HfO2 CoO-CoCr2O4-Cr2O3-HfO2  necessitated cross-section preparations. Per alloy two cross-sectional SEM/BSE micrographs taken at two magnifications (left: general view, and right: detailed view) are presented in Fig. 8 for the two 0.25C-containing alloys and in Fig. 9 for the three 0.50C-containing alloys. First, one can globally see that the total thickness of the external scales (including all types of oxides) tends to be higher for the Hf-rich alloys than for the Hf-free ones. This is in good agreement with the mass gain hierarchies.  9  Illustration of the surface states of the three 0.50C-containing alloys after oxidation test (cross-sectional SEM/BSE examin ations, SEM/EDS oxide determinations);  top:  the Hf-free Co2 alloy, middle:  the Hf-rich CoHf2a alloy, bottom:  the Hf-rich  CoHf2b alloy  416  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  \\x0c', 'Be r thod and Con ra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  Table 6  Contents in chromium and in hafnium in outermost surface (average, standard deviation and range determined from values obtained in five randomly chosen locations); EDS spot analysis  Elements in outermost surface  Average wt-%Cr  Std deviation wt-%Cr  wt-% Hf range  Co1 CoHf1 Co2 CoHf2a CoHf2b  13.1 8.2 13.9 13.0 17.6  4.1 2.2 1.7 0.5 1.2  / 0.0 / 0.0 0-0.04  Second, instead of being the main oxide existing in the scales, chromia is less present than CoO and than the spinel oxide. Chromia is generally situated in the most inner position, inside the alloy where it constitutes a more or less continuous and geometrically complex internal layer. Just above one meets the spinel oxide and after, the most external one, the cobalt oxide CoO. It seems that the external cobalt oxide CoO layer has grown outward by cationic diffusion of Co2+ while the medium  spinel oxide CoCr2O4 and the internal chromia Cr2O3 both have grown inward by anionic diffusion of O2−. This scheme seems validated by the position of the hafnium oxides HfO2 which resulted from the in situ oxidation of the HfC carbides: they are disseminated inside chromia and the spinel oxide but never in the outer cobalt oxide scale. The oxidised HfC carbides seem playing the role of the gold marks that one sometimes deposits on slightly pre-oxidised alloys in order to reveal the cationic or anionic character of diffusion during the subsequent high temperature oxidation. One can underline that, concerning the observed diffusion phenomena, the studied alloys did not differ from what is commonly known about more common cobalt-based alloys8 and notably ternary Co-Cr-C and ternary Co-Cr alloys. The presence of hafnium did not influence the diffusion phenomena usually observed for the existing cobalt-based alloys, which is not surprising since Hf is exclusively present as HfC carbides (no Hf atoms in solid solution in the matrix), carbides which are in situ transformed into isolated internal HfO2 oxides without any diffusion of the Hf atoms.  10  Cr and Hf concentration profiles from the {oxide scale/alloy} interface (left side) toward the bulk (right side), perpendicularly  to the sample surface (profiles plotted from EDS spot analyses performed perpendicularly to the {external oxide scale/alloy}  interface on increasingly deep locations in the alloys)  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  417  \\x0c', 'Be r thod and Conra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  The alloys’ chemical compositions in outermost surface were specified by EDS spot analysis, five times per oxidised alloys. The results, presented in Table 6, reveal that the chromium contents have become rather low: around 13 wt-% for three of the alloys, and even 8 wt-% for the CoHf1 alloy. These values are not compatible with the sustainability of any chromia-forming behaviour, which was effectively lost as demonstrated by the analyses of the present oxides (major part is composed of cobalt oxide and of spinel oxide). The Co1, Co2 and CoHf2a alloys (13-14 wt-%Cr in outermost surface) present a rather regular/linear scale/alloy interface. In contrast oxidation has clearly become catastrophic for the CoHf1 and CoHf2b alloy (mass gain of about 0.01 g cm−2 after 50 h at 1100°C). This resulted in a chromium content of only 8 wt-%Cr in outermost surface (very low) for the first alloy, and in a Cr content of almost 18 wt-% for the second one, a particularly high value which shows that the alloy started to be consumed by the inward progression of the catastrophic oxidation. Concerning Hf it is almost absent in the outermost surface as in the matrix elsewhere. Obviously the HfC carbides close to the scale/ alloy interface are transformed in oxides in situ without any release of Hf atoms. Indeed the EDS profiles given in Fig. 10 show that the Hf content is often zero and jumps at high level when HfC carbides are touched by the electron beam of the SEM. On these profiles one can see that the chromium content regularly but rapidly decreases from a 50 to 80 µm depth in the bulk when one goes to the scale/alloy interface where it reaches very low levels. In the case of the CoHf2b alloy for which catastrophic oxidation penetrated the alloy the Cr content at the oxidation front is maintained at a rather high level because of the oxidation of all the elements belonging to the alloy.  Conclusions  Hafnium carbides present important advantages for strengthening cast superalloys: their script-like morphology, their imbrication with matrix in an eutectic compound, their very high stability at high temperature, notably. This, associated with an intrinsically mechanically resistant cobalt-based matrix, recently proved that they allowed high mechanical properties at levels of working temperature considered forbidden for polycrystalline cast alloys. However, it was here demonstrated that the same combination of a cobalt matrix reputed to be less resistant against high temperature oxidation than corresponding nickel alloys (essentially for Cr diffusion easiness reasons), with HfC carbides made from a particularly oxidable element, does not allow oxidation resistance high enough to envisage industrialisation and use in the present state. Further work needs to be carried out. The improvement of these alloys in oxidation behaviour requires additional work and changes in the chemical composition, for example by enriching them in chromium beyond the actual 25 wt-% (30 wt-% or more) or by addition of aluminium and/or silicon. Cr deposition by pack-cementation or by other chemical vapour deposition techniques can be also considered. In another  418  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  side, the high easiness of hafnium oxidation is probably not the single reason of the deterioration of the oxidation resistance of the base alloys Co-25Cr-0.25 and 0.50C. The cross-sectional observations clearly showed that the HfC carbides initially present close to the oxide/alloy interface did not disappear unlike chromium carbides or tantalum carbides for example. They are transformed into HfO2 oxides which stay on place even close to the oxidation front and then obstruct chromium diffusion along the grain/interdendritic boundaries. This is possible that they permanently disturb the supply of the oxidation front in chromium and thus favour a precocious generalised oxidation by loss of the chromia-forming behaviour: outward cationic Cr3+ diffusion shortly replaced by inward anionic O2− diffusion. The effect of HfC on the different phenomena governing the oxidation process remains to be analysed and clearly explained.  Acknowledgements  The authors would like XRD experiments.  thank Pascal Villeger  for  the  Funding  This work is totally free (no funding, no grant).  References  5.  6.  7.  3.  (eds.):  ‘The  1972,  superalloys’,  1. C. T. Sims and W. C. Hagel New York, John Wiley & Sons. 2. C. T. Sims, N. S. Stoloff and W. C. Hagel (eds.): ‘Superalloy II - high temperature materials for aerospace and industrial power’, 1987, New York, John Wiley. P. Berthod, S. Michon, L. Aranda, S. Mathieu and J. C. Gachon: ‘Experimental and thermodynamic study of the microstructure evolution in cobalt-base superalloys at high temperature’, Calphad, 2003, 27, 353-359. 4. M. J. Donachie and S. J. Donachie (eds.): ‘Superalloys - a technical guide’, 2nd edn, 2002, Materials Park, OH, ASM International. P. Berthod, S. Michon, L. Aranda and P. Steinmetz: ‘Consequences of the high temperature microstructure evolution on the tensile properties at high temperature of a cobalt-based superalloys reinforced tantalum carbides’, ‘Matériaux 2006’, Dijon, France, by Proc. November 2006. P. Berthod, C. Heil and L. Aranda: ‘Influence of the morphologic evolution of the eutectic carbides at high temperature on the thermal expansion behaviour of refractory cast alloys’, J. Alloys Compd., 2010, 504, 243-250. P. Berthod: ‘High temperature properties of several chromium-containing Co-based alloys reinforced by different types of MC carbides (M=Ta, Nb, Hf and/or Zr)’, J. Alloys Compd., 2009, 481, 746-754. 8. D. J. Young: ‘High temperature oxidation and corrosion of metals’, 2008, Amsterdam, Elsevier. 9. W. S. Lee and G. M. Kim: ‘The effect of Hf, Y and Zr additions on the oxidation behaviour of Ni-based Inconel 601 at high temperature’, Han’guk Pusik Hakhoechi, 1995, 24, 124-133. 10. K. Ishii, M. Kohno, S. Ishikawa and S. Satoh: ‘Effect of Ti, Zr and Hf on high-temperature oxidation resistance of Fe-20Cr-5Al-0.1La alloy foils’, Mater. Trans. JIM, 1998, 39, 1040-1045. 11. X. Chunmei, J. Guo and F. Yang: ‘High-temperature oxidation behaviour of NiAl-28Cr-5Mo-1Hf alloy’, Jinshu Xuebao, 2001, 37, 857-860. 12. H. Guo, L. Sun, H. Li and S. Gong: ‘High temperature oxidation behaviour of hafnium modified NiAl bond coat in EB-PVD thermal barrier coating system’, Thin Solid Films, 2008, 516, 5732- 5735.  \\x0c', 'Be r thod and Con ra th  Ox ida t ion a t 1100 °C o f H fC -s t reng thened Co a l loys  13.  14.  P. Berthod and E. Conrath: ‘Mechanical and chemical properties at high temperature of {M-25Cr}-based alloys containing hafnium carbides (M=Co, Ni or Fe): creep behavior and oxidation at 1200°C’, J. Mater. Sci. Technol. Res., 2014, 1, (1), 7-14. P. Berthod: ‘Kinetics of high temperature oxidation and chromia volatilization for a binary Ni-Cr alloy’, Oxid. Met., 2005, 64, 235-252.  15.  16.  P. Berthod: ‘Oxidation start detection on heating parts of thermogravimetry curves for high temperature alloys based on nickel, cobalt or iron’, Open Corr. J., 2011, 4, 1-8. P. Berthod: ‘Thermogravimetric study of oxide spallation for chromium-rich cast cobalt-based and iron based alloys oxidized at high temperature’, Open Corr. J., 2009, 2, 61-70.  Canadian Metallurgical Quarterly  2016  VOL 55  NO 4  419  \\x0c']"
},{
  "_id": 157,
  "PDF": "Oxidation behaviour of a pressureless sintered HfB2-MoSi2 composite.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 29 (2009) 1809-1815  Oxidation behaviour of a pressureless sintered HfB2-MoSi2 composite  ∗  Diletta Sciti  , Andrea Balbo, Alida Bellosi  CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy  Received 28 July 2008; received in revised form 22 September 2008; accepted 24 September 2008  Available online 5 November 2008  Abstract  The thermal stability of a 80-vol.% HfB2 + 20 vol.% MoSi2 composite is tested under oxidizing environment. Oxidation tests are carried out in temperatures ≥1200 flowing synthetic air in a TG equipment from 1000 to 1400 C with exposure time of 30 h. At C the silica resulting from oxidation of molybdenum disilicide seals the sample surface, preventing hafnium diboride from fast degradation. Analysis of the kinetics is carried out through fitting of the thermogravimetric curves. Between 1200 and 1400 C, the kinetic curves deviate from a parabolic behaviour, being more close to a logarithmic-parabolic behaviour. © 2008 Elsevier Ltd. All rights reserved.           Keywords: Ceramic; Microstructure; Oxidation; Borides  1.  Introduction  Hafnium diboride is gaining increasing attention as a highly refractory compound with superior properties at high temperature. Nowadays the design and production of new materials suitable to withstand high temperatures are stimulated by an increasing demand for applications in the field of thermal protection systems and for several industrial sectors like foundry. HfB2 -based composites display a number of unique properties including hardness, high thermal and electrical conductivity and chemical stability.1-12 HfB2 -SiC composites are able to withstand very high temperatures5-7 and can maintain their room C in air.6,7 The addition of temperature strength up to 1500 MoSi2 to borides seems to be very promising as it allows the densification of composites by pressureless sintering and the resulting composites retain their strength up to 1500 HfB2 -MoSi2 composites have also been tested in an arc jet facility at 1900 C showing excellent stability due to the development of a silica-based scale.12 So far, oxidation studies reported in the literature have concerned pure HfB2 ,13,14 HfB2-SiC5,8 materials, and HfB2-Si3N4 15 materials. Oxidation data for pure HfB2 between 1200 and 1700 C could be fitted to a parabolic rate equation. Around 1700 C there was an abrupt increase of  C.10,11                 ∗  Corresponding author. Tel.: +39 0546699748 fax: +39 054646381.  E-mail address: diletta.sciti@istec.cnr.it (D. Sciti).  0955-2219/$ - see front matter © 2008 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2008.09.018        the oxidation rate due to transition from monoclinic to tetragonal HfO2 .13 More recently, a mechanistic model that interprets the oxidation behaviour of Zr and Hf borides in the range 1000-1800 C has been formulated.14 At temperature below 1400 C, the rate limiting step is the diffusion of dissolved oxygen through liquid boria, while at higher temperatures boria is lost by evaporation and the oxidation rate is limited by diffusion of molecular oxygen between columnar blocks of HfO2 .14Silica has very low diffusivities for oxygen and is a very protective scale at moderate temperatures. Hence, SiC, Si3N4 and transition metal silicides are considered very useful additives for improving the oxidation resistance in the middle temperature range. HfB2 -SiC materials5,8 were reported to withstand temperatures as high as 1700 C thanks to the formation of a borosilicate glass containing HfO2 crystals, which is an effective barrier against the inward diffusion of oxygen. For HfB2-Si3N4 materials oxidation studies were carried out15 up to 1400 C. At temperatures in the 900-1200 C range, the main oxidation products were HfO2 , B2O3 and borosilicate glass. In the 1200-1400 C range, SiO2 and HfSiO4 were detected. Kinetic curves recorded at 1450 and 1600 C were shown to be parabolic.In this work, the oxidation behaviour of a HfB2 -20 vol.% MoSi2 composite is tested in the middle temperature range, i.e. between 1000 and 1400 C. Actually, MoSi2 was already found to be beneficial for improving the oxidation resistance of ZrB2-20 vol.% MoSi2 composites in the same temperature range.16-18                    \\x0c', '1810  D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815  Fig. 1.  (a) Polished microstructure and (b) X-ray diffraction of as-sintered HfB2 -20 vol.% MoSi2 .  2. Experimental procedure  2.1. Material  Commercial powders were used for the preparation of the composite material: HfB2 (Cerac Incorporated, Milwaukee, USA), particle size range 0.5-5 \\u242em. Impurities: Al (0.07%), Fe (0.01%), Zr (0.47%); MoSi2 (<2 \\u242em, Aldrich, Milwaukee, USA), mean particle size 2.8 \\u242em and oxygen content of about 1 wt.%. Dense pellets were prepared by pressureless sintering at 1950 C. Details on the material’s processing are reported elsewhere.11     2.2. Oxidation tests  sized 9.0 mm × 8.0 mm × 1.0 mm were Rectangular plates cut from the sintered pellet. The specimens were cleaned in ultrasonicated acetone bath, dried and weighed (accuracy 0.01 mg). The oxidation tests were carried out in a thermogravimetric analyser (model STA449, NETSCH, Geraetebau GmbH, Selb, Germany), in synthetic air (composition: 80 vol.% N2 + 20 vol.%  Fig. 2. Thermogravimetric curves of the HfB2 -20 vol.% MoSi2 composite in the range 1000-1400 C.           O2 , with 30 ml/min gas flow) between 1000 to 1400 C with isothermal exposure time of 30 h for each experiment, heating rate 30 C/min and free cooling. The fast heating-up stage prior to the isothermal period was applied to minimize oxidation effects before reaching the target temperatures. The mass varia−3 mg of sensitivity. The tion was recorded continuously with 10 TG measurements evaluation was performed with the subtraction of Buoyancy effect corrections. As-sintered and oxidized sample surfaces were analysed by X-ray diffraction (Cu K␣ radiation, Miniflex Rigaku, Tokyo, Japan). Surfaces and polished cross-sections were analysed by scanning electron microscope (SEM, Leica Cambridge S360, Cambridge, UK) and energy dispersive microanalysis (EDS, Model INCA energy 300; Oxford Instruments, High Wycombe, UK).  3. Results  3.1. Microstructure of the as-sintered material  The sintered material contained a low amount of residual porosity, 2%, as ascertained by SEM analysis and density measurements. Crystalline HfB2 , MoSi2 and traces of HfO2 were identified by X-ray diffraction (Fig. 1a). An example of the polished section is displayed in Fig. 1b. HfB2 grains have a rounded shape while the MoSi2 phase has a very irregular morphology with very low dihedral angles. This peculiar characteristic indicates that MoSi2 was very ductile at the sintering temperature or could have formed a liquid phase, which is not surprising since the sintering temperature was close to its melting point (2020 C). The analysis of secondary phases by EDS confirmed the presence of HfO2 , HfC and traces of a Mo-B phase. Further details are reported elsewhere.11     3.2. Oxidation curves  Thermogravimetric curves recorded during the oxidation of the composite 80 vol.% HfB2 + 20 vol.% MoSi2 are displayed in Fig. 2. With the exception of the curve collected at 1100 C showing a mass loss, in all of the cases a weight gain was recorded. The weight gain after oxidation at 1200 C was        \\x0c', 'D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815  1811     lower than after oxidation at 1000 C. It was found that none of the recorded kinetics followed a simple law, i.e. linear or parabolic. Hence, kinetic curves were fitted according to a multiple law model including a linear, parabolic and logarithmic contribution18 :  \\x01W  S  = Kpar · t 1/2 + Klin · t + Klog · log(t )  (1)  where \\x01W/S is the weight gain per unit surface, Kpar , Klin , Klog are the parabolic, linear and logarithmic term, respectively. The physical meaning of Eq. (1) is that different phenomena can simultaneously take place during oxidation, such as diffusion of inward and outward species, rupture of the scale, crystallization  of new phases inside the scale, evaporation of scale material, etc.18 Ceramic composites are highly unlikely to show a simple oxidation behaviour, thus a multiple regression analysis seems to be more suitable for the interpretation of experimental data. With the exception of the linear term (Klin ) which can be negative when significant development of gaseous products occurs, negative K parameters do not reflect a physical reality. The experimental data and the fitting curves are displayed in Fig. 3a-f. Kinetic parameters and fit goodness values, R2 , are reported in Table 1. The oxidation kinetics recorded at temperatures of 1000 and 1100 C cannot be properly fitted by Eq. (1) either because the R2 values (fit goodness) are too low or because K parameters assume negative values. The most     Fig. 3.  (a-f) Fitting of the thermogravimetric curves in the range 1000-1400     C. The model is the sum of the linear, parabolic, logarithmic contributions.  \\x0c', '1812  D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815     C)  1000  Oxidation temperature (  Table 1 Kinetic parameters of the oxidation fitting curve, relatively to linear (Klin ), parabolic (Kpar ) and logarithmic (Klog ) contributions and R2 (fit goodness) values. −2 min −1 ) −2 min Klin (mg cm Kpar (mg cm 6.6 × 10 −2 × 10 −8 × 10 −8 × 10 1 × 10 -  −5 −5 −5 −5 −4  1200 (t > 800 min)  1200 (0 < t < 800)  −1/2 )  0.003  0.015  1100  1300  1350  Klog  0.49  -  -  0.16  0.09  -  -  -  1400  -  0.022  0.12  R2  0.905  0.960  0.975  0.994  0.988  0.997  0.992           reasonable solution is a linear behaviour starting after a transition stage, as shown in Fig. 3a and b. During oxidation at 1000 C a weight gain is observed, while at 1100 C the negative linear term indicates that an overall mass loss occurs, i.e. reactions leading to mass loss prevail on reactions leading to mass gain. The interpretation of the kinetic curve at 1200 C is quite difficult. The fit with Eq. (1) produces a negative logarithmic term which is not physically acceptable. The data seem to indicate that two subsequent stages occur during oxidation. The fit with Eq. (1) however does not allow for a shift in time after which a new law acts exclusively, but interprets the data in terms of different kinetics acting at the same time. Thus, in Fig. 3c the experimental data are interpreted in terms of a linear stage occurring for 100 < t < 800 min and a logarithmic stage occurring for t > 800 min. The curve recorded at 1300 C was successfully fitted by Eq. (1), as displayed in Fig. 3d. The plot shows the separate contributions of the linear, parabolic and logarithmic term and their sum, which is overlapped to the experimental data. As can be seen, the 1300 C kinetics is dominated by the logarithmic term. At 1350 and 1400 C, the data can be fitted by Eq. (1), but the linear term assumes negative values. Assuming that at these temperatures no significant loss of gaseous species occurs due to formation of a stable oxide, a logarithmic-parabolic curve is the best fitting solution for the experimental data (Fig. 3e and f).           3.3. Oxidation product        The crystalline phases formed after oxidation are shown in Fig. 4. The morphological evolution of surfaces and crosssections is reported in Figs. 5a-d and 6a-d. For samples oxidized in the temperature range 1000-1200 C, monoclinic HfO2 was the main crystalline product due to oxidation. In the 1300-1400 C oxidation range, beside HfB2 from the bulk, HfO2 and HfSiO4 were detected. A slight reduction of HfB2 peaks intensity was observed upon oxidation at 1400 C, due to thickening of the surface scale, as illustrated later. In the investigated temperature range, all the sample surfaces were covered by a continuous borosilicatic glassy layer, in which small crystals (HfO2 /HfSiO4 ) were dispersed (Fig. 5). This layer had a thickness of less than a few micrometers and completely sealed the sample surface. The presence of boron was clearly              detected by EDS after oxidation at 1300-1400 C, (see inset in Fig. 5c). Increasing wrinkling of the surface oxide was observed with increasing the oxidation temperature, due to formation of gaseous products. The polished cross-sections (Fig. 6) give further indications. C, a 1-\\u242em thick silica-based layer Upon oxidation at 1000 forms. The coverage is irregular and hafnium oxide crystals are observed. Upon oxidation at 1100 and 1200 C, the surface scale has a thickness of about 2-4 \\u242em. In the 1300-1400 C range, formation of bubbles on the surface becomes more evident and, despite the presence of the surface oxide, subsurface oxidation takes place. The surface borosilicate glass has a very irregular thickness of about 5-10 \\u242em, irrespective of the oxidation temperature. The subsurface modified layer has a thickness which increases from 10 to 20 \\u242em increasing the oxidation temperature from 1300 to 1400 C. In this intermediate layer penetration of the glassy phase and formation of molybdenum oxide (MoO) or MoB or a mixed MoOB phase is observed, as shown in Fig. 6. The morphology of these phases suggests solidification from a liquid phase.        4. Discussion  In this paragraph a correlation between TG curves, XRD data and SEM analyses is attempted. The oxidation reactions occur Fig.  4. X-ray  diffraction  of  the  sample  after  oxidation  in  the  range  1000-1400  C.     \\x0c', 'D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815  1813  Fig. 5. SEM analysis of oxidized surfaces at: (a) 1100     C, (b) 1200     C, (c) 1300     C and (d) 1400     C. Inset: EDS spectrum of the surface.  ring in the HfB2-MoSi2 composite during high-temperature treatments in air depend on the oxidation reactions of the two constituent phases and on further interactions among their oxide products. HfB2 oxidizes according to the following reactions13 : HfB2 (s) + (5/2)O2 (g) = HfO2 (s) + B2O3 (l) = B2O3  B2O3  (g).  (l),  (1)  (2)  At temperature >1100 C, liquid boria starts to vaporize.13 The oxidized surface is constituted of a hafnia-based layer. MoSi2 has an excellent oxidation resistance at temperatures C because of a protective silica surface layer17 : >1000 (s) + (7/2) (g) = 2SiO2 (s) + MoO3  MoSi2  O2     (s).     (3)     In the subsurface layer, at T ≥ 1300 C, liquid boria from reaction (1) can further react with MoSi2 , to form MoB or MoOB or MoO species, in agreement with SEM observations. In the composite of the present work, according to microstructural observations, the oxidation during isothermal runs in the tested temperature range can be explained as follows.  •     1000 C: The oxidation of HfB2 into HfO2 according to reactions (1 and 2) and formation of silica from reaction (3) are the dominant effects. Liquid B2O3 is known to have a protective effect, while HfO2 is semi-protective due to its anion        deficiency.13 Due to the low thickness of the oxidized layer, loss of volatile species cannot be ruled out. The roughly linear behaviour suggests that the scale is not fully protective. 1100 C: The negative linear kinetic curve indicates that loss of volatile species prevails. The evaporation of liquid boria is thought to play a major role at this temperature (reaction (2)). The morphology of the oxidized layer and X-ray diffraction data indicate the formation of a glassy layer embedding hafnia crystals. The glassy layer, however, is probably too thin to exert an efficient protective action. 1200 C: As previously observed, the weight gain at the end of the isothermal stage is lower than at 1000 and 1100 C. At this temperature, evaporation of liquid boria should be the dominant mechanism but SEM analyses demonstrate the formation of a continuous surface glassy layer. During the first stage controlled by linear kinetics (for 100 min <t < 800 min) the formed scale could be too thin to effectively protect the sample from mass loss due to evaporation of liquid boria, hence weight gain and weight loss phenomena occur simultaneously. During the second stage (t > 800 min), the formation of a stable silica-based scale starts to limit gas evolution and hence the kinetics has a decelerating feature, being the superimposition of a negative linear term and a positive logarithmic term. 1300-1400 C: In this temperature range, the major contribution to the weight gain comes from the formation of a        •  •  •  \\x0c', '1814  D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815  Fig. 6. SEM analysis of oxidized sections at: (a) 1100     C, (b) 1200     C, (c) 1300     C and (d) 1400     C.  composite, which was tested in oxidizing environment using the same experimental conditions of the present work.20 Compared to the ZrB2 -based system, the weight gain after 30 h of isothermal run is notably reduced. At 1200 C it is about 1/2 of the ZrB2 weight gain, at 1300 C 1/4, at 1400 C 1/5. A direct comparison of the weight gain after isothermal stage at 1400 C is                 SiO2 -rich glass that increases in thickness with increase of the oxidation temperature. Furthermore, we have observed the formation of hafnon crystals. Since HfO2 crystals are embedded in a continuous and partially amorphous silica-rich layer, further reaction can occur among the oxidation products to form crystalline HfSiO4 . Similar to the mechanism formation of zircon,20 when HfO2 grains are embedded of in a silica layer, interstitial silicon diffuses and dissolves into crystalline hafnia until the solubility limit is reached, thereafter HfSiO4 crystals precipitate. The decelerated kinetics at 1300 C contains a strong logarithmic term. Logarithmic oxidation is thought to set in when the oxide begins to crystallize because the decrease of the amorphous phase leads to a decrease of the effective oxidation cross-section, thus reducing the oxidation rate constant.19,21 In the composite analysed, the logarithmic term could account for the crystallization of hafnon from HfO2 and glassy silica. The kinetics recorded at 1350 and 1400 C are still far from a parabolic behaviour, but it can be observed that the logarithmic component tends to decrease and the parabolic component to increase. The parabolic term can be explained in terms of diffusion of oxygen through the glassy layer.18 The logarithmic contribution derives from the crystallization of hafnon in the glassy scale.     Oxidation data collected for the HfB2-20 vol.% MoSi2 composite, can be compared with data from a ZrB2 -20 vol.% MoSi2  Fig.  7. Comparison  of weight  gain  behaviour  between ZrB2 -MoSi2  and  HfB2 -MoSi2 composites.  \\x0c', 'D. Sciti et al. / Journal of the European Ceramic Society 29 (2009) 1809-1815  1815     shown in Fig. 7. Since the secondary phase type and amount is the same for the two composites, it must be concluded that the higher stability of the HfB2 -composite is related to the oxidation resistance of the matrix. This feature was already found in the study of monolithic ZrB2 and HfB2 .13 Besides, for both systems, the weight gain at 1200 C is lower than at 1000 C due to formation of the protective silica scale. Other similarities concern the formation of MoB in the subsurface layer and the formation of zircon (hafnon) in the silica scale at 1300-1400 C. However, differently from the logarithmic curves displayed by the HfB2 -MoSi2 system, the kinetics of the ZrB2-MoSi2 system are mostly parabolic in the 1200-1400 C temperature range, indicating that the dominant mechanism is diffusion of oxygen through the silica-based oxide layer and that other phenomena such as crystallization of the scale play a minor role.           5. Conclusions        Oxidation tests carried out in a thermogravimetric analyser proved that the composite 80 vol.% HfB2 + 20 vol.% MoSi2 can withstand temperatures up to 1400 C under oxygen-containing environment. At lower temperatures (1000-1200 C), Hafnium oxide is the main oxidation product. At temperatures ≥1200 C, owing to the oxidation of MoSi2 , appreciable amounts of glass form, which improves the oxidation resistance of the material. The kinetic curves generally deviate from a parabolic behaviour, being more close to a logarithmic-parabolic behaviour, which is the result of the complex reactions occurring in the scale. Compared to the previously analysed ZrB2 -20 vol.% MoSi2 , the HfB2 -based composite shows a higher stability.     Acknowledgments  The authors wish to thank L. Silvestroni for preparation of the bulk material and S. Guicciardi for useful discussion on the kinetic curves.  References  3. Wuchina, E., Opeka, M., Causey, S., Buesking, K., Spain, J., Cull, A. et  al., Designing for ultrahigh-temperature applications: the mechanical and thermal properties of HfB2 , HfCx , HfNx and ␣Hf(N). J. Mater. Sci., 2004, 39, 5939-5949.  4. Opeka, M. M., Talmy, I. G. and Zaykoski, J. A., Oxidation-based materials     selection for 2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J. Mater. Sci., 2004, 39, 5925-5937.  5. Gasch, M., Ellerby, D., Irby, E., Beckman, S., Gusman, M. and Johnson,  S., Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics. J. Mater. Sci., 2004, 39, 5925-5937.  6. Bellosi, A., Monteverde, F. D. and Sciti, D., Fast densification of ultra high-temperature ceramics by spark plasma sintering. Int. J. Appl. Ceram. Technol., 2006, 3, 32-40.  7. Monteverde, F., Ultra-high temperature HfB2 -SiC ceramics consolidated by hot-pressing and spark plasma sintering. J. Alloy Compd., 2007, 428,  197-205.  8. Monteverde, F. and Bellosi, A., The resistance to oxidation of an HfB2 -SiC composite. J. Eur. Ceram. Soc., 2005, 25, 1025-1031.  9. Monteverde, F., Hot pressing of hafnium diboride aided by different sinter additives. J. Mater. Sci., 2008, 43, 1002-1007.  10. Sciti, D., Silvestroni, L.  and Bellosi, A., Fabrication and properties of  HfB2 -MoSi2 composites produced by hot pressing and spark plasma sintering. J. Mater. Res., 2006, 21, 1460-1466.  11. Silvestroni, L. and Sciti, D., Effects of MoSi2 additions on the properties of Hf- and Zr-B2 composites produced by pressureless sintering. Scripta Mater., 2007, 57, 165-168.  12. Savino, R., De Stefano Fumo, M., Silvestroni, L. and Sciti, D., Arc-jet testing  on HfB2 and HfC-based ultra-high temperature ceramic materials. J. Euro. Cer. Soc., 2008, 28, 1899-1907.  13. Berkowitz-Mattuck, J. B., High-temperature oxidation, III zirconium and hafnium diborides. J. Electrochem. Soc., 1966, 113, 908-914.  14. Parthasarathy, T. A., Rapp, R. A., Opeka, M. and Kerans, R. J., A model for the oxidation of ZrB2 , HfB2 and TiB2 . Acta Mater., 2007, 55, 5999-6010. 15. Klein, R., Desmaison-Brut, M., Desmaison, J., Mazerolles, L. and Trichet,  M. F. In High Temperature Corrosion and Protection of Materials 6, Parts  1 and 2, Proceedings 2004, vol. 461-464, pp. 849-856.  16.  Jeng, Y. L. and Lavernia, E. J., Review: processing of molybdenum disilicide. J. Mater. Sci., 1994, 29, 2557-2571.  17. Natesan, K. and Deevi, S. C., Oxidation behaviour of molybdenum silicides and their composites. Intermetallics, 2000, 8, 1147-1158.  18. Sciti, D., Brach, M. and Bellosi, A., Oxidation behaviour of a pressureless sintered ZrB2 -MoSi2 ceramic composite. J. Mater. Res., 2005, 20(4), 922-930.  19. Nickel, K. G., Multiple law modeling for the oxidation of advanced ceramics  and a model-independent figure of merit. In Corrosion of Advanced Ceram ics, ed. K. G. Nickel. Kluwer Academic Publishers, Norwell, MA, 1994, p.  59.  1. Opeka, M. M., Talmy, I. G., Wuchina, E. J., Zaykoski, J. A. and Causey, S.  J., Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds. J. Eur. Ceram. Soc., 1999, 19, 2405-2414.  2. Fahrenholtz, W. G., Hilmas, G. E., Talmy, I. G. and Zaykoski, J. A., Refractory diborides of zirconium and hafnium. J. Am. Ceram. Soc., 2007, 90,  20. Veytizou, C., Quinson, J. F., Valfort, O. and Thomas, G., Zircon formation  from amorphous silica and tetragonal zirconia: kinetic study and modelling.  Solid State Ionics, 2001, 139, 315.  21. Ogbuji, L. and Singh, M., High temperature oxidation behaviour of reactionformed silicon carbide ceramics. J. Mater. Res., 1995, 10(12), 3232-3240.  1347-1364.  \\x0c']"
},{
  "_id": 158,
  "PDF": "Oxidation behaviour of coarse and fine SiC reinforced ZrB2 at re-entry and atmospheric oxygen pressures.pdf",
  "Text": "['Ceramics International 46 (2020) 11056-11065  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Oxidation behaviour of coarse and fine SiC reinforced ZrB2 at re-entry and atmospheric oxygen pressures  T  Rubia Hassana,1, Rishabh Kundub,1, Kantesh Balania,∗  a Department of Materials Science and Engineering, Indian Institute of Technology, Kanpur, Kanpur, 208016, b Department of Ceramic Engineering, National Institute of Technology, Rourkela, Odisha, 769008, India  India  A R T I C L E  I N F O  A B S T R A C T  (coarse ~30 μm and fine ~5 μm) were ZrB2 composites reinforced with SiC of two different particle sizes processed using spark plasma sintering. The oxidation behaviour of the synthesised composites was investigated at 1500 °C under two different partial pressures of oxygen. At higher oxygen partial pressure of 0.21 atm, coarse SiC reinforced ZrB2 composite developed thicker oxide layers as compared to that of fine SiC reinforced ZrB2 composite. The thicker silica layer provides an enhanced oxidation resistance for extended exposure periods of greater than 3 h. However, at lower oxygen partial pressure of 2 × 10−5 atm, fine SiC reinforced ZrB2 composite showed better oxidation resistance due to segregation of finer ZrO2 particles on the surface, which prevented inward transportation of oxygen.  Keywords: ZrB2 Oxidation Silica Oxygen partial pressure SiC depletion layer  1.  Introduction  Ultra-High-Temperature Ceramics (UHTCs) is a family of materials having melting temperatures > 3000 °C and potential of using at high temperatures [1]. Ceramic non-oxide compounds of refractory borides, nitrides and carbides, such as HfB2, ZrB2, HfN, TaC etc. are few examples that belong to the aforementioned family of materials. Research interest in UHTCs has heightened in the past few years because of their potential to produce components for hypersonic aerospace vehicles, such as thermal protection systems, gas turbine parts, cutting tools, refractory linings etc. [2-7]. UHTC composites based on borides of Zr and Hf, constitute a class of favoring materials for use in extreme applications such as control surface components (sharp leading/trailing edges) of hypersonic vehicles, because of their high melting temperatures (> 3000 °C), high thermal conductivity (> 50 W/mK) , better oxidation resistance than that of other members of the UHTC family, dimensional stability, and retained strength at elevated temperatures [3,8-13]. However, as an aerospace material, ZrB2 has superiority over HfB2, being lighter while possessing proportionate levels of oxidation resistance [8,9,11,14-17]. Resistance to oxidation is a primary concern for the application of UHTCs. Enormous efforts have been dedicated to this afore parameter [18-22]. Due to the low oxidation resistance and strength/fracture toughness, the physical and mechanical properties of monolithic ZrB2 are not appropriate for the above specific application.  Low sinterability of ZrB2 due to strong covalent bonding and low selfdiffusion coefficient add to its limitations [23]. However, a modern sintering technique called spark plasma sintering (SPS) has offered several advantages over other conventional sintering techniques like hot pressing and pressure-less sintering, in the processing of UHTCs. The application of electric field coupled with externally applied pressure offers the benefit of achieving fast heating rate, low sintering low dwell time with nearly fully dense compact [24-26]. temperature, Second phase (e.g. SiC, MoSi2) introduction has been successful in significantly improving the oxidation resistance as well as mechanical properties of UHTCs. It has been reported that high SiC addition is favourable for oxidation resistance at high temperatures (up to 1800 °C) in air due to the formation of silica glass. However, it is found that at a reduced partial pressure of oxygen, this trend gets reversed, due to rapid active oxidation of SiC, which gives rise to the formation of negligible amount of silica glass [23,27]. Detailed oxidation studies (1500 °C, in air) quantitatively report the inhomogeneous oxidation behaviour of ZrB2-20 vol.% SiC composite. The oxidation product that is the outermost borosilicate-glass layer (whose thickness varies with the oxidation conditions) is closely involved in the inhomogeneous the ZrB2-SiC composite [28]. The oxidation of oxidation behaviour of the ceramic composites is reported to be diffusion controlled, and the activation energy of oxygen in the oxide scale of the ceramics oxidised under temperatures below or at 1500 °C is 5.1 kJ/mol for ZrB2-20 vol.%  ∗ Corresponding author. E-mail address: kbalani@iitk.ac.in (K. Balani). 1 Have contributed equally as first author.  https://doi.org/10.1016/j.ceramint.2020.01.125 Received 5 September 2019; Received in revised form 31 December 2019; Accepted 14 January 2020  Available online 17 January 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  SiC UHTC [29]. It has also been reported that the oxidation resistance can be improved by; controlling the starting size of the powder and the additives, incorporating efficient techniques for mixing, and using production techniques that minimise the secondary phase formation [30]. In accordance with the previously reported results, a diboride matrix composite having SiC only as its second phase, behaves as one of the most promising composition. M. Mashhadi et al. [31] studied the effect of SiC content and particle size on the sintering behaviour and mechanical properties of ZrB2. The relative density was higher for nanosized SiC as compared to micrometer-sized SiC reinforced ZrB2 composites and increased up to 15 wt % for micrometer-sized SiC and up to 10 wt % for nano-sized SiC reinforced composites. Hardness followed the same trend both with the SiC content and SiC particle size. The effect of SiC particle on the oxidation behavior of ZrB2 has not been studied yet though the effect of SiC content has been investigated by J. Han et al. [27], as discussed above. The purpose of this study is, thus, to investigate the oxidation resistance of ZrB2 containing 20 vol.% SiC of two different particle sizes as the oxidation behaviour on the optimisation of the microstructure has not been extensively studied at higher temperatures.  2. Material and methods  2.1. Preparation and characterization of materials  ZrB2 commercial powder was obtained from H.C. Starck, Germany (average particle size of 8.2 μm, 99% + pure), coarse SiC commercial powder which was obtained from Samics Research Materials, Bareilly, India (average particle size of d50 = 86.3 μm, 90% pure with imUP, purities of SiO2 and Al2O3, see Ref. [21]) and fine SiC commercial powder was obtained from H.C. Starck, Germany (average particle size of d50 = 5 μm, 99% pure). Fig. 1 shows the micrographs, and corresponding histograms used for particle size measurement of the starting powders. Fig. 1 depicts the D10, D50 and D90 of the starting powders. The powders were weighed in suitable proportions. To prepare the powder mixtures of ZrB2 + 20 vol.% coarse SiC (ZSC) and ZrB2 + 20 vol.% fine SiC (ZSF) were then dry ball milled using WC balls (ball: powder ratio of 5:2) for 8 min at 500 RPM (with the particle size distribution presented in Appendix I). The milled powder was then spark plasma sintered (SPS, Dr. Sinter 515S, Japan) with a heating rate of 100 °C/min and at an applied pressure of 30 MPa for 10 min at a maximum temperature of 1850 °C under vacuum, to prepare pellets of 15 mm diameter and 5 mm thickness. These pellets were, then, ground using 15 μm diamond slurry and polished using abrasive papers till a mirror-like surface was formed. Theoretical density and bulk density were obtained using the ruleof-mixture (by taking the density of ZrB2 as 6.08 g/cm3 and 3.21 g/cm3 for SiC) and the Archimedes method, respectively. Phases were identified by X-ray diffraction (XRD, Rigaku MiniFlex 600) operated at 25 kV, and 15 mA with Cu-Kα (λ = 1.541 Å) radiation scanned at a step size of 0.02° and a scan speed of 0.16°/s. The microstructures of the samples were examined using secondary electron and backscattered electron modes of imaging in a scanning electron microscope (SEM, WSEM, JSM-6010LA, JEOL, Japan), while as energy dispersive spectroscopy (EDS) was used to determine the chemical compositions of different phases.  2.2. Oxidation tests  Sample specimen of 10 × 3.6 × 5 mm3 dimensions from ZSC as well as ZSF were sectioned using a slow-speed diamond saw followed by polishing. Samples were successively ultrasonically cleaned in deionised water and ethanol and dried before exposure. Samples were loaded on an alumina crucible and then exposed to oxygen at partial pressures of 0.21 and 2 × 10−5 atm for 180 min at 1500 °C, using a tube furnace (OKAY Furnace, Kolkata, India) with zirconia heating  element. The gas composition in the furnace was maintained at 21 vol. % oxygen, and 79 vol.% argon to obtain 0.21 atm oxygen partial pressure condition, whereas furnace was maintained at 99.99 vol.% argon (assuming balance 0.01 vol.% of impurity as air) to obtain the oxygen partial pressure of 2 × 10−5 atm. Phases after oxidation were characterised using X-Ray diffraction and Raman spectroscopy (Princeton Instruments, STR Raman, TE-PMT detector) using Nd-YAG green laser source (λ = 532 nm) for an acquisition time of 4 seconds per accumulation (20 accumulations) at 50% of 30 mW laser power. Composition and microstructure of the cross-section and the surface of the samples after oxidation were obtained by EDS and SEM.  3. Results and discussion  3.1. Physical and microstructural characterisation  The relative densities of ZSC and ZSF obtained by Archimedes’ principle are 89.0% and 98.6%, respectively, as depicted in Table 1. The relative density increased as the SiC grain size changed from coarse (D50 = 86.3 μm) to fine (D50 = 5 μm), which is in accordance with results reported by Ref. [32]. The starting SiC particle size plays a vital role in controlling the final microstructure of the composite, which is evident through observations tabulated in Table 1. Fig. 2(a) and (b) shows SEM micrographs of the polished surfaces of ZSC and ZSF, respectively. The decrease in particle size of SiC coarse after sintering is because of the breakdown of the initial powder particles and a negligible grain growth during sintering (8 min milling time did not reduce the size of ZrB2 and fine SiC particles). Though the size of coare SiC has changed from D50 of 86.3 μm to D50 of 31 μm. The SEM micrographs of milled powders is provided in Appendix I in the supplementary information and SiC coarse particle size (D50) obtained from the Figure (b) of Appendix I is reported in Table 1. The SiC phase, which is indicated by dark grey features in Fig. 2(a) and (b), is dispersed homogeneously within the ZrB2 matrix. Grain growth of ZrB2 was better inhibited due to the presence of fine SiC particles than coarse SiC particles as depicted in Table 1. SiC reinforcement has substantially enhanced densification of ZrB2 during SPS by promoting liquid phase formation during SPS, leading to an increased densification [33]. It may be pointed out that the 10% impurity in coarse-SiC is Al2O3 and SiO2, and as only 20 vol% SiC is added to ZrB2, which brings down the total oxide impurity level to 2%. Further, the oxides either break down or get reduced after SPS as noted in XRD of the SPS pellets [21] indicating that oxide impurities are mostly removed. On the other hand, fine-SiC is ~99% + pure. So, in summary, the impurity content in the two sizes of SiC reinforcement may not significantly alter the oxidation behaviour of SPS processed composites. The representative scanning electron micrographs of the oxidised surfaces of ZSC and ZSF at 1500 °C under 0.21 atm oxygen partial pressure (the normal oxygen partial pressure in the troposphere, the last layer encountered by aerospace vehicles during re-entry) are shown in Fig. 3(a) and (b), respectively. A glassy silica layer appears to have covered both the surfaces. Passive oxidation of SiC forms silica as per Equation (1) which, in turn, leads to the formation of the glassy silica layer. Aggregated white crystals of various shapes and sizes formed amidst the glassy silica layer in both the samples as indicated in the insets of Fig. 3(a) and (b). Energy-Dispersive Spectroscopy (EDS) analysis verified these as zirconia (Equation (3)). The appearance of zirconia on the surface was most probably because of the formation and the coalescence of the gaseous products (i.e., CO, SiO) formed as a result of oxidation of SiC as depicted in Equation (2) which probably gave rise to bubbles that facilitated the transportation of zirconia (which formed due to the oxidation of ZrB2 present in the matrix as per Equation (3)) to the glassy silica layer [27]. When the vapour pressure exceeds the ambient pressure, the bubbles burst to led to the deposition of zirconia crystals inside the externally forming glassy silica layer. The formation of the zirconia particles amidst the glassy silica layer is also  11057  \\x0c', \"R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Fig. 1. SEM micrographs taken in SE mode showing morphology of as-procured powders of (a) ZrB2; (b) SiC coarse [21], and (c) SiC fine, and histograms for grain size measurement of (d) ZrB2; (e) SiC coarse, and (f) SiC fine.  Table 1  Composition and Density of the samples.  Sample ID  Composition  Particle size (D50) before milling  Particle size (D50) after milling  Grain size after sintering  Density (g/cm3)  % Theoretical Densification (ρA/ρT)  ZrB2 (μm)  SiC (μm)  ZrB2 (μm)  SiC (μm)  ZrB2 (μm)  SiC (μm)  ρT  ρG  ρA  ZSC  ZSF  ZrB2 + 20 vol.% coarse SiC ZrB2 + 20 vol.% fine SiC  8.2  8.2  86.3  5.0  8.2  8.2  31  5.0  9.0 ± 3.8  32.5 ± 11.4  5.51  4.91  4.91  89.0  7.0 ± 2.2  6.8 ± 1.2  5.51  5.11  5.43  98.6  ρT: Theoretical density; ρG: Geometrical density; ρA: Archimedes' density.  in accordance with the binary phase diagram of ZrO2-SiO2 system where at 1500 °C, ZrSiO4 (solid solution of silica and zirconia as stated in the phase diagram) and SiO2 phases are present [34]. Passive oxidation of SiC (> 1100 °C) [35].  Active oxidation of SiC (> 1400 °C) [35].  SiC (s)  +  O (g)  2  SiO (g)  CO (g)  +  =  Oxidation of ZrB2 (> 1100 °C) [36].  SiC (s)  +  3 2  O (g)  2  SiO (g)  2  CO (g)  +  =  ZrB (s)  2  +  5 2  (1)  O (g)  2  ZrO (s)  2  =  +  11058  B O (g)  2  3  (2)  (3)  \\x0c\", 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Fig. 2. SEM micrographs taken in back scattered imaging mode of the polished surfaces of (a) ZSC and (b) ZSF along with corresponding EDS spectra.  The scanning electron micrographs of the oxidised surfaces of ZSC and ZSF at 1500 °C under 2 × 10−5 atm oxygen partial pressure in the mesosphere (an approximate 1/10,000 of mean-sea-level oxygen partial pressure, encountered by aerospace vehicles during re-entry) are shown in Fig. 4(a) and (b), respectively. The surface of ZSC was covered by  segregated particles of silica and zirconia, whereas the surface of ZSF was covered by segregated particles of zirconia with a negligible amount of silica entrapped amidst the zirconia particles. The active oxidation of SiC (Equation (2)) takes place at a significantly lower oxygen partial pressure of 2 × 10−5 atm [37], due to which SiO  Fig. 3. SEM (secondary electron imaging mode) micrographs with zoomed insets and EDS spectra for the oxidised (a) ZSC and (b) ZSF at 1500 °C and 0.21 atm oxygen partial pressure.  11059  \\x0c', 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Fig. 4. SEM micrographs (secondary electron imaging mode) and EDS spectra for the oxidised (a) ZSC and (b) ZSF at 1500 °C and 2 × 10−5 atm oxygen partial pressure.  Fig. 5. Cross-sectional micrographs (back scattered imaging mode) of oxidised materials at 1500 °C for 3 h: (a) ZSC, pO2 = 0.21 atm; (b) ZSF, pO2 = 0.21 atm; (c) ZSC, pO2 = 2 × 10−5 atm and (d) ZSF, pO2 = 2 × 10−5 atm. (Bright and dark features in unaffected region of all samples represent ZrB2 and SiC respectively. Bright and dark (darker and rougher) features in SiC depleted region of all samples represent ZrB2 and SiC respectively, wherein in SiC is partly affected by oxygen. Bright (diffused) and dark (diffused) features in ZrO2+SiO2 region of all samples represent ZrO2 and SiO2, respectively. Bright elongated features in dark SiO2 rich layer (in (a)) represent ZrO2.  11060  \\x0c', \"R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Table 2  Thickness of  Composite  the layers of ZrB2-SiC oxidised at 1500 °C under different oxygen partial pressures compared with the literature.  Td (μm)  To (μm)  Ts (μm)  Comments (w.r.t. oxidation resistance)  Reference  pO2 = 0.21 atm ZrB2-20SiC 40 To + Ts = 300 ZSC 117 74 ZSF 44 34 pO2 = 2 × 10−4 - 2 × 10−5 atm ZrB2-20SiC 85 To + Ts = 310 ZSC 370 40 ZSF 54 160 (12 + 148)  9-94 4  ~0 ~0  Higher (To + Ts) and lesser Td → Preferred at higher pO2. Delayed formation of a thick Ts → Preferred at higher pO2 for extended exposure period. Instantaneous formation of a thin Ts → Preferred at higher pO2 for short exposure period.  J. Han et al. [27] Current work Current work  High Td → Preferred at higher pO2. Higher Td and coarser ZrO2 → Preferred at higher pO2 for extended exposure period. Lesser Td and finer segregated ZrO2 → Preferred at lower pO2.  J. Han et al. [27] Current work Current work  Tt: Total Thickness of the layers; Td: Thickness of the SiC-depleted layer; To: Thickness of the oxide layer (ZrO2-SiO2); Ts: Thickness of the silica layer.  higher than the vapour pressure of CO (oxidised product of SiC). The scanning electron micrographs of the oxidised cross-sections of ZSC and ZSF at 1500 °C under 0.21 atm oxygen partial pressure are shown in Fig. 5(a) and (b), respectively. A general observation in both the cross-section micrographs presented four distinct layers: (i) a silicarich outer layer, which is believed to act as a protective scale and formed due to passive oxidation SiC as per Equation (1); (ii) a subscale of crystalline zirconia (Equation III), containing silica; (iii) a zirconium diboride region with a minute amount of SiC with altered topography ‘‘SiC-depleted region’‘), which (the is believed to form due to incomplete surface oxidation of SiC grains at the experimental temperature and period of exposure used; and (iv) unaltered or the unoxidised material which has been earlier observed by J. Han et al. [27]. On one hand, the glassy silica layer observed in ZSC is non-homogenous, but contrastingly that is not the case for ZSF in our current work. The thickness of the various aforementioned layers of both the oxidised samples is summarised in Table 2. The table also compares the results obtained with results reported by J. Han et al. [27]. All the aforementioned layers of ZSF were thinner than that of ZSC's, which resulted in ZSF's total layer thickness to be less than that of ZSC's. The scanning electron micrographs of the oxidised cross-sections of ZSC and ZSF at 1500 °C under 2 × 10−5 atm oxygen partial pressure are shown in Fig. 5(c) and (d), respectively. The micrograph of ZSC showed three distinct layers: (i) an upper layer of crystalline zirconia, containing silica (Equations (1) and (3)); (ii) a zirconium diboride region with a minute amount of SiC with altered topography (the ‘‘SiCdepleted region’‘), which is believed to happen due to incomplete surface oxidation of SiC grains at the experimental temperature and period of exposure used; and (iii) unaltered or the unoxidised material. The micrograph of ZSF also showed the same three distinct layers, but there was a repetition of the crystalline zirconia-silica layer after the “SiC depleted region”. The thickness of each layer of both the oxidised samples is summed up in Table 2. The SiC depleted region was thinner in ZSF whereas the zirconia-silica layer was thinner in ZSC, but the total layer thickness of ZSC was greater than that of ZSF. The reoccurrence of the crystalline zirconia-silica layer below the “SiC depleted region” is because the SiO2 that formed due to passive oxidation of SiC as per Equation (1) and the ZrO2 formed as per Equation (3) did not get sufficient energy to diffuse out to the surface because of low exposure time and thus remained entrapped. The total thickness of the layers in ZSC, as well as ZSF, increased with a decreasing oxygen partial pressure, which signifies that at lower oxygen partial pressure, rapid oxidation takes place. Low oxygen partial pressure also facilitated active oxidation of SiC to SiO and CO, which is indicated by the increase in the layer thickness of the SiC depleted region at low oxygen partial pressure. The active oxidation of SiC led to formatiation of negligible silica formation, wherein the glassy silica layer disappeared for both ZSC and ZSF with a decreased oxygen partial pressure and favoured inward transportation of oxygen, resulting in an increased oxidation of the composites. Similar observations have been reported by J. Han et al. [27]. Due to higher porosity in ZSC (~11%) than ZSF (~1.4%), at higher  Fig. 6. XRD pattern of SPSed pellets.  Fig. 7. XRD pattern of oxidised materials at 1500 °C.  becomes the primary oxidised product formed instead of SiO2, and hence a glassy silica layer was not observed instead silica particles were observed due to minor passive oxidation of SiC as per Equation (1). Zirconia, formed as a result of oxidation of ZrB2, is present in the matrix as per Equation (3). ZSF's surface also contained C (as detected by EDS shown in Fig. 4(b)), which may have been deposited due to the high partial pressure of Argon in the experimental environment which was  11061  \\x0c\", 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Fig. 8. Raman spectra of oxidised samples (a) ZSC, pO2 = 0.21 atm; (b) ZSF, pO2 = 0.21 atm; (c) ZSC, pO2 = 2 × 10−5 atm and (d) ZSF, pO2 = 2 × 10−5 atm.  Table 3  Raman peaks matched with the literature along with their assigned modes.  Peak  ZrO2  ZrO2/SiO2  Si SiC  Raman shift (cm−1) observed  Raman shift (cm−1) from literature  Assigned mode  Reference  141 172.5-177 215 324-332 343-347 510 565 642 481-490  520 784-788 971-975  138-148 177-179 221-222 330-333 349 502-504 558-566 637-640 471-475/480-488  521 789-804 965-972  Bg Ag Bg Ag Ag Bg Ag Ag Ag/Si-O-Si bending TO TO LO  [38-43]  [38,44-46]  [40,47] [48,49]  oxygen partial pressure, the initial penetration of oxygen (because of an easier path through the pores) was more, which yielded layers in ZSC with thickness more than that of ZSF. Due to enhanced penetration, swift and non-uniform oxidation of SiC took place which formed a thicker glassy silica layer (as per Equation (1)) as well as “SiC depleted region” in ZSC. The non-homogenous silica layer also formed because of the different oxidation rates of SiC in various parts of the material. Due to negligible porosity in ZSF, the same phenomenon did not occur, and, thus, fast and non-uniform oxidation of SiC did not take place, and hence it resulted a thinner glassy silica layer as well as “SiC depleted region” than ZSC. For prolonged periods, ZSC would be more resistant  to oxidation as compared to ZSF due to the thicker glassy silica layer (formed at an early stage), which would prevent the further inward transportation of oxygen. ZSC has an inferior degree of interconnectivity among the SiC grains than ZSF because the finer SiC grains in ZSF (Fig. 2(b)) enhanced the liquid phase sintering during SPS as a result of which SiC got distributed uniformly and continuously along the grain boundaries of the ZrB2 matrix. Due to the higher degree of interconnectivity of SiC in ZSF, it should have been less resistant to oxidation as reported by J. Han et al. [27], though that is not the case observed at the lower oxygen partial pressure condition as the finer segregated zirconia (0.7 ± 0.3 μm) on the surface of ZSF than ZSC (formed zirconia size: 5.6 ± 1.3 μm) trapped the liquid phase in the zirconia layer and prevented oxygen transport through the liquid as stated previously by E. Eakins [30], thereby checking oxidation of the SiC and the ZrB2 matrix. It can, thus, be concluded that the effect of the entrapment of the liquid phase by finer segregated zirconia is greater than the effect of degree of interconnectivity among the SiC grains for the oxidation of the composites.  3.2. Phase analysis  Fig. 6 shows the XRD pattern of the unoxidised ZSC as well as ZSF, where the characteristic peaks of ZrB2 and SiC are clearly marked, and no secondary phases have been detected. Fig. 7 shows the XRD pattern of ZSC, as well as ZSF, oxidised at 1500 °C under 0.21 atm oxygen partial pressure. The primary phases observed were silica and zirconia, whose characteristic peaks are clearly marked. ZrO2 phase, as expected below SiO2, is captured by XRD pattern. Few peaks of very low intensity of SiC are also observed. Fig. 7 also shows the XRD pattern of ZSC, as  11062  \\x0c', 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  Fig. 9. Representation of the surface and cross sectional features of oxidised materials at 1500 °C for 3 h: (a) ZSC, pO2 = 0.21 atm; (b) ZSF, pO2 = 0.21 atm; (c) ZSC, pO2 = 2 × 10−5 atm and (d) ZSF, pO2 = 2 × 10−5 atm.  well as ZSF, oxidised at 1500 °C under 2 × 10−5 atm oxygen partial pressure. The primary phases detected in ZSC were ZrO2 and SiO2, whereas zirconia was the only phase present in ZSF, as observed in Fig. 5(c) and (d) and represented in Fig. 8. The characteristic peaks are clearly marked. The XRD patterns verify the phases formed at the surfaces of the composites after oxidation as estimated from SEM micrographs and EDS spectra. Fig. 8 compares the features exhibited by the materials after oxidation at 1500 °C for 3 h. The Raman spectra collected from the cross-sections of all the oxidised samples are presented in Fig. 8. Spectra were collected from each sample at different locations labelled in the spectra, within different regions (as marked in the SEM micrographs). The thickness of different layers measured from SEM micrographs was taken into account while collecting the spectra so that a characteristic spectrum from each region of interest is obtained and compared with the others. The peaks obtained in the oxidised samples are compared in Table 3. ZrO2 peaks were obtained in SiO2 rich layer of both coarse and fine SiC reinforced samples at pO2 = 0.21 atm (Fig. 8(a) and (b)), which extend down to SiO2 + ZrO2 region in all the four samples (Fig. 8(a), include ~141 cm−1 (b), Fig. 8(c) and (d)). These peaks in ZSF at pO2 = 2 × 10−5 atm, ~170 cm−1 peak in ZSC at pO2 = 0.21 atm and ZSF at pO2 = 2 × 10−5 atm, ~214 cm−1 peak and ~350 cm−1 peak in ZSC and ZSF at pO2 = 0.21 atm and ZSC at pO2 = 2 × 10−5 atm, cm−1 ~324-332 peak in ZSC at pO2 = 0.21 atm and ZSF at pO2 = 2 × 10−5 atm, ~481.5-491 cm−1 peak in all the samples corresponds to either SiO2 or ZrO2 or both as strong ZrO2 peak usually appears at 471-475 cm−1 and SiO2 peak appears at 480-488 cm−1, cm−1 510 peak in ZSC at pO2 = 0.21 atm and ZSC at pO2 = 2 × 10−5 atm, ~565 cm−1 peak in ZSC at pO2 = 0.21 atm and ~640 cm−1 peak in ZSC and ZSF at pO2 = 2 × 10−5 atm [38,39,41,46]. The peak at ~520 cm−1 in ZSC at pO2 = 0.21 atm and ZSC at pO2 = 2 × 10−5 atm corresponds to the transverse optical phonon (TO band) of crystalline Si which could have formed due to the reduction of SiC into elemental Si [40]. A sharp peak was also found in the XRD pattern of ZSC at pO2 = 0.21 atm and a small peak in ZSC at pO2 = 2 × 10−5 atm (2θ ~ 34.5°) corresponding to elemental Si.  Raman peak at ~788 cm−1and ~971 cm−1 in the un-altered material correspond to the transverse optical phonon (TO band) and longitudinal optical phonon (LO band) scattering of SiC [48,49] which confirms that the SiC did not undergo oxidation in this region. The Raman peaks obtained in different layers of all the samples are in confirmation with the SEM micrographs (Fig. 5). Fig. 9 schematically compares the features exhibited by the materials after oxidation at 1500 °C for 3 h. At pO2 = 0.21 atm, ZSC (Fig. 9(a)) formed an outer silica-rich layer with an underneath SiO2+ZrO2 layer followed by SiC depletion region. At same pO2 of 0.21 atm, ZSF (Fig. 9(b)) also developed the same layers in similar order with each layer being thinner than that of ZSC. At pO2 = 2 × 10−5 atm, the outer most silica-rich layer was not found in ZSC (Fig. 9(c)) due to active oxidation of SiC into SiO. As a result, the ZrO2 rich layer was followed by SiC depletion region. However, the limited passive oxidation of SiC into SiO2 yielded some SiO2 particles on the surface. Finer zirconia particles formed at the surface of ZSF oxidised at pO2 = 2 × 10−5 atm (Fig. 9(d)), helps entrap oxygen better than the coarser zirconia particles formed at the surface of ZSC (Fig. 9(c)), thereby, keeping oxidation under check to a greater extent. ZSC is expected to exhibit better oxidation resistance at pO2 = 0.21 atm for extended periods of greater than 3 h due to thicker silica layer formation than ZSF, which is initially inhomogeneous, but for shorter exposure periods, ZSF is expected to possess better oxidation resistance due to the instantaneous formation of the silica layer.  4. Conclusions  ZrB2 + 20% vol. SiC composites having coarse SiC (D50 = 31 μm) and fine SiC (D50 = 5 μm) particle sizes were processed using SPS. ZrB2 + 20% vol. SiC composite reinforced with coarse SiC exhibited an initial low resistance to oxidation at 0.21 atm oxygen partial pressure due to a greater extent of transportation of oxygen through pores, but would be more resistant to oxidation for longer exposure periods than ZrB2 + 20% vol. SiC composite with fine SiC as the former contains a thicker glassy silica layer. At an oxygen partial pressure of 2 × 10−5 atm, ZrB2  11063  \\x0c', 'R. Hassan, et al.  Ceramics International 46 (2020) 11056-11065  + 20% vol. SiC composite having fine SiC exhibited better oxidation resistance than its counterpart and may continue to do so even for longer exposure periods at the lower oxygen partial pressure condition due to the segregation of finer zirconia on the surface. Finer SiC grain size is, therefore, preferable for the oxidation resistance of the material at high oxygen partial pressure (for short exposure periods) and low oxygen partial pressure (for short as well as long exposure periods). The present study is focused on temperature ~1500 °C, which is typically the exposure temperature experienced during the re-entry conditions of aerospace vehicles. Finer SiC grain size also inhibits the exaggerated growth of ZrB2 particles during sintering as compared to that of coarse SiC, but coarse SiC grain size is preferred in the composite for better oxidation resistance for longer exposure periods at high oxygen partial pressure. Therefore, a necessary trade-off has to be made between them concerning the actual requirement of the application conditions.  Author contribution  Rubia Hassan designed the experiments. Rubia Hassan and Rishabh Kundu performed the experiments, analysed the data and wrote the paper. Kantesh Balani conceived the idea, obtained the funding, and guided in organising and editing the paper.  Declaration of competing interest  Authors declare no competing interest for the results reported in this manuscript.  Acknowledgements  Rubia Hassan at Indian Institute of Technology Kanpur (IITK) acknowledges the MHRD Govt. of India for funding through Prime Minister Research Fellowship. KB acknowledges IMPRINT funding (IMP/2018/000739) from Science & Engineering Research Board (SERB), Department of Science & Technology, Govt. of India. Advanced Centre of Materials Science (ACMS) at IITK for extending characterisation facilities (SEM and XRD) is acknowledged. Electro-ceramics Research lab under Dr Shobit Omar at IITK for extending the facility of environment-controlled furnace is also acknowledged.  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2020.01.125 .  References  [1]  F. Monteverde, C. Melandri, S. Guicciardi, Microstructure and mechanical properties of an HfB2 + 30 vol.% SiC composite consolidated by spark plasma sintering, Mater. Chem. Phys. 100 (2006) 513-519. [2] D. M. Van Wie, D. G. Drewry, D. E. King, and C. M. Hudson, \"The Hypersonic Environment: Required Operating Conditions and Design Challenges,\" pp 59155924.. [3] M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, \"Oxidation-based Materials Selection for 2000°C + Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience,\" pp. 5887-5904. [4] M. Shahedi Asl, B. Nayebi, M.G. Kakroudi, M. 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},{
  "_id": 159,
  "PDF": "Oxidation behaviour of HfB2-15 vol._ TaSi2 at low, intermediate and high temperatures.pdf",
  "Text": "['Available online at www.sciencedirect.com  Scripta Materialia 63 (2010) 601-604  www.elsevier.com/locate/scriptamat  Oxidation behaviour of HfB2-15 vol.% TaSi2 at intermediate and high temperatures  low,  Diletta Sciti,* Valentina Medri and Laura Silvestroni  CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy  Received 9 March 2010; revised 25 May 2010; accepted 29 May 2010  Available online 4 June 2010  The thermal stability of a HfB2-15 vol.% TaSi2 composite was tested by thermogravimetric analysis from 25 to 1500 °C, in a bottom-up furnace at 1600 °C for 15 min and in an arc-jet facility at temperatures between 2000 and 2500 °C. The layered scales resulting from arc-jet and conventional furnace oxidation are very similar. The flexural strength does not change after oxidation at 1600 °C. The effect of TaSi2 on the oxidation resistance is investigated.  Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: Ceramic; Microstructure; Oxidation; Borides  Hafnium diboride (HfB2) is gaining increasing attention as a highly refractory compound able to withstand aggressive environments such as those demanded for thermal protection systems [1-10]. Oxidation studies reported in the literature concern pure HfB2 [10,11], HfB2-SiC [4,5], HfB2-Si3N4 [6] and HfB2-MoSi2 materials [8,9]. SiC, Si3N4 and transition metal silicides are considered useful additives for improving the oxidation (1000-1700 °C) resistance at intermediate temperatures due to the low diffusivity of oxygen through silica. HfB2-SiC materials [4,5] were reported to withstand temperatures as high as 1700 °C thanks to the formation of a borosilicate glass containing HfO2 crystals. HfB2- MoSi2 composites tested in an arc-jet facility at 1900 °C showed excellent stability due to the development of a stable silica-based scale [9]. In this paper, oxidation tests are conducted on a HfB2-based composite doped with 15 vol.% TaSi2. This material has very interesting high-temperature properties, such as a flexural strength of 600 MPa at 1500 °C, which suggests good high-temperature stability [12]. Addition of TaSi2 has been reported to significantly improve the oxidation resistance of ZrB2-SiC composites, due to an increased viscosity of the glass diffusion barrier [13-16] or to oxidation and breakup of TaB2-ZrB2 boride solid-solution grains into fine particles which retain the liquid phase [17]. Hence the purpose of this pa * Corresponding author. Tel.: +39 0546699748;  fax: +39 054646381;  e-mail: diletta.sciti@istec.cnr.it  is  per is to determine whether the addition of TaSi2 beneficial for the oxidation behaviour of HfB2. The composite 85 vol.% HfB2-15 vol.% TaSi2 was produced from commercial powders: HfB2 (Cerac Incorporated, Milwaukee, USA), \\x00325 mesh, d50 = 1 lm; TaSi2 (ABCR, GmbH & Co, Karlsruhe, Germany), \\x0045 lm. Details on powder processing are reported elsewhere [12]. The powder mixture was consolidated by hot-pressing, at 1900 °C/15 min and with a pressure of 30 MPa, achieving a final relative density of 99% [12]. Oxidation tests were carried out on rectangular plates 9.0 \\x02 8.0 \\x02 1.0 mm3 in a thermogravimetric (TG) analyzer (model STA449, NETSCH, Geraetebau GmbH, Selb, Germany), in synthetic air (80 vol.% N2 + 20 vol.% O2, with 30 ml min-1 gas flow) between 25 and 1500 °C, using a heating rate of 10 °C min-1 and free cooling. The oxidation tests in the bottom-up loading 1600 °C 13 \\x02 2.5 \\x02 2 mm3 bars in static air. Specimens were lofurnace was conducted at for 15 min on cated in the furnace when the maximum temperature was reached and then removed and air quenched after the exposure time was complete. For arc-jet tests, hemispheric and conical models were exposed to a mixture 20 vol.% O2-80 vol.% N2 at a mass flow rates of up to 1 g s-1. The models were located at a distance of 1 cm from the torch exit, and the chamber working pressure was between 100 and 500 Pa. The average specific total enthalpy in the proximity of the specimens was evaluated to be 10 and 16 MJ kg-1 for the hemisphere and the cone, respectively. The surface temperature was  1359-6462/$ see front matter Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2010.05.050  \\x0c', '602  D. Sciti et al. / Scripta Materialia 63 (2010) 601-604  solution with composition (Hf0.85Ta0.15)B2. The formation of this peculiar structure was disclosed to be due to complex phenomena occurring during sintering, including TaSi2 decomposition, formation of liquid phases and interdiffusion of metal cations [12]. Not all of the starting TaSi2 formed solid solutions; indeed, in the sintered material its content was 7 vol.%, as estimated from XRD spectra [12]. The non-isothermal TG plot up to 1500 °C is reported in Figure 2a. A weight increase was recorded starting from 800 °C. Around 1200 °C, a slight attenuation of mass gain was observed. At 1300 °C the weight started to increase again up to 1400 °C, after which a marked weight loss occurred. The final weight gain per unit surface was around 1 mg cm-2. After completion of the TG analysis, XRD revealed the formation of monoclinic HfO2, the principal reflection of hexagonal B2O3, and traces of Ta2O5 and Hf6Ta2O17. The surface of the oxidized samples is constituted by a silica-based glass and isolated regions where HfO2 grains and deposits of B2O3 were recognized by EDS (Fig. 2b,c). The cross-section (Fig. 2d) displayed a surface glassy layer and an inner oxidized layer constituted by a borosilicate glass and Ta-doped HfO2 grains. After oxidation at 1600 °C/15 min in the bottom-up furnace, the main crystalline phase was a mixed oxide with stoichiometry Hf6Ta2O17 and traces of m-HfO2. The sample surface was covered by an amorphous silica-based scale, with embedded platelet-like crystals (Fig. 3a). The cross-section analysis revealed that the scale was a multilayered oxide with approximate thickness of 100 lm, adhering well to the bulk. According to EDS analysis, the outer layer (I) comprised an amorphous silica layer, containing 1.6 at.% Hf and 2.5 at.% Ta (spectrum A in Fig. 3c), and Hf6Ta2O17 crystals (spectrum B in Fig. 3c). Underneath this layer, the scale (layer II) mainly consisted of of large, round HfO2  Figure 1. HfB2-15 vol.% TaSi2  composite:  (a) polished section;  (b)  TEM image showing the core-shell structure.  measured through a radiation ratio pyrometer (Infratherm ISQ5, Impac Electronic Gmbh, Germany). The as-sintered material and oxidized specimens were examined on the surface using X-ray diffraction (XRD; Siemens D500, Germany) to identify the crystalline phases. Microstructural modifications induced by oxidation were analyzed by scanning electron microscopy (SEM; Cambridge S360) and energy dispersive spectroscopy (EDS; INCA Energy 300, Oxford instruments, UK). On the sintered material, local phase analysis was performed using transmission electron microscopy (TEM; FEI, CM12, Eindhoven, The Netherlands) equipped with an energy-dispersive X-ray system (EDS, EDAX Genesis 2000, Ametek GmbH, Wiesbaden, Germany). After oxidation at 1600 °C, the retained strength of the oxidized specimens was compared to that of the as-sintered ones in a universal screw-type testing machine (Instron 6025) in three point-bending tests on chamfered bars 13 \\x02 2.5 \\x02 2 mm3. Three specimens for oxidized and as-sintered materials were tested. A fine microstructure with little porosity (<1%) was observed in the as-sintered material (Fig. 1a). HfB2 grains have a core-shell structure (Fig. 1b), where the core is the original HfB2 phase, whilst the shell is a solid  Figure 2.  (a) Weight gain per unit surface (S) from 25 to 1500 °C; (b) sample surface after completion of the TG run; (c) detail showing deposits of  B2O3;  (d) cross-section.  Figure 3. 1600 °C Oxidation. (a) Surface morphology. (b) Layered cross-section: I, Hf6Ta2O17 crystals in Ta-borosilicate glass; II, (Hf,Ta)O2 crystals after grain growth and partially SiOx-filled porosity; III, fine-grained HfO2. (c) EDX spectra of the regions indicated in (b).  \\x0c', 'D. Sciti et al. / Scripta Materialia 63 (2010) 601-604  603  Figure 4. Arc jet: (a) hemisphere surface morphology (inset: Hf6Ta2O17 crystals in Ta-borosilicate glass); (b) cross-section showing the layered scale;  (c) cone surface morphology (inset: Hf6Ta2O17 crystals in Ta-borosilicate glass and large cavities); (d) cross-section of oxide spallation.  (\\x185 lm) containing <3 at.% Ta (spectrum C in grains Fig. 3c) and porosity. Finally, the inner layer (III) contained fine HfO2 grains (<2 lm) and residual porosity. The flexural strength of the as-sintered material was 768 ± 26 MPa. After oxidation the strength was 731 ± 20 MPa, i.e. not significantly changed. Concerning the arc-jet tests, the hemisphere underwent three consecutive runs during which the maximum temperature reached about 2000 °C on the tip. The total time of exposure was 11 min. The altered morphology of the model is shown in Figure 4a, b. The sample surface displayed an amorphous silica-based scale with embedded Hf6Ta2O17 platelet-like crystals and a regular pattern of pores (see inset in Fig. 4a). Along the crosssection the scale appeared as a multilayered oxide of variable thickness (130 lm in the near-tip region, 50 lm in the back), composed of the above-described layers I-III as for the sample treated in the furnace at 1600 °C (Fig. 4b). Sporadic formation of macrocracks at the oxide/bulk interface during cooling suggests possible problems of adherence to the bulk. No silica was observed in the intermediate and inner layers that exhibited some porosity. It was observed that layer I is about 1/3 of the entire scale thickness as for the sample treated in the bottom-up furnace. The conical specimen underwent two consecutive runs during which the maximum temperature ranged between 2400 and 2500 °C on the tip, with a total time of exposure of 8 min. Indeed, the analysis of the sample surface revealed that the extent of damage was much more pronounced for the conical model, due to the higher temperature reached, especially in the near-tip regions. During cutting and polishing operations, detachment of a portion of the tip scale occurred, indicating that this part of the scale was poorly bonded to the material bulk due to mismatch of the coefficients of thermal expansion (Fig. 4c). The surface morphology was heavily altered by the shear forces associated with the hot stream, which enhanced the bursting of bubbles and the formation of craters (see inset in Fig. 4c). The cross-section analysis confirmed that the scale was a layered oxide of variable thickness (\\x18170 lm near the tip, \\x18100 lm in the back), but in this case the distinction between layers I-III was not so obvious, due to severe damage. When present, it was found that the thickness of layer I was much reduced, to less than 1/5, compared to the overall scale thickness. Underneath, the scale mainly comprised large, round HfO2 grains (5 lm) containing <3 at.% of Ta (layer III) and large porosities. Macrocracks and spallation were observed at the interface between oxide and unreacted bulk (Fig. 4d).  known that HfB2 oxidizes  It is well [10,11]: HfB2 ðsÞ þ 5=2O2 ðgÞ ¼ HfO2 ðsÞ þ B2O3 ðlÞ  B2O3 ðlÞ ¼ B2O3 ðgÞ  according  to  ð1Þ  ð2Þ  Boron oxide is partially protective at temperatures <1000 °C but has a low melting point and high vapour pressure, and therefore starts to evaporate at T > 1100 °C, reducing the effectiveness of the diffusion barrier. According to recent studies [11], at temperatures below 1400 °C, the rate-limiting step is the diffusion of dissolved oxygen through liquid boria, while at higher temperatures boria is lost by evaporation and the oxidation rate is limited by diffusion of molecular oxygen along columnar blocks of HfO2. On the other hand, TaSi2 may oxidize according to: 2TaSi2 þ 6:5O2 ðgÞ ¼ Ta2O5 þ 4SiO2  ð3Þ  Reaction (3) is favourable over the entire temperature range. According to available phase diagrams, Ta2O5 and B2O3 form a eutectic point at 1700 °C [18]. Hence it is reasonable to hypothesize that Ta2O5, SiO2 and B2O3 may form liquid phases at temperatures around 1700 °C or lower. Another peculiar finding of this work is the formation of the mixed oxide with stoichiometry Hf6Ta2O17. Based on the JCPDF database, Hf6Ta2O17 is the only ternary crystalline phase in the Hf-Ta-O system. Zhao et al. studied HfTaxOy layers deposited on SiO2 and deduced that this phase crystallized from an amorphous Hf,Ta-rich silica matrix at 950 °C [19]. In the conditions of the TG experiments, i.e. non-isothermal run between room temperature and 1500 °C, the response to oxidation of this composite is not very different from other borides containing SiC or Si-containing phases. Reactions (1) and (2) become active around 800 °C. In particular, as the boride grains have a core- rim substructure it is likely that reaction (1) for the rim becomes:  ðHf 0:85Ta0:15 ÞB2 ðsÞ þ 5=2O2 ðgÞ ¼ ðHf 0:85Ta0:15 ÞO2 ðsÞ þ B2O3 ðlÞ  ð4Þ  The attenuation of weight gain observed around 1200 °C indicates that below this temperature the production of silica through reaction (3) is not fast enough to offer protection against escape of gaseous species, which according to previous studies on the oxidation kinetics is the rate-limiting step between 1000 and 1200 °C [5-8]. At T P 1300 °C, a consistent borosilicate layer starts to form [5-8]. Above 1400 °C weight loss prevails over  \\x0c', '604  D. Sciti et al. / Scripta Materialia 63 (2010) 601-604  weight gain. The surface features, consisting of isolated regions of Ta-doped HfO2 grains embedded in the liquid and surface deposits of B2O3, suggest that the weight loss could be mostly due to boria evaporation. The observed characteristics resemble those displayed by ZrB2-SiC composites after oxidation at 1500 °C, as reported by Karlsdottir et al. [20]. In analogy with their findings, the surface arrangement might be a consequence of convective movements of a fluid boria-rich liquid where some Ta-doped hafnia is dissolved. In the same regions, the residual B2O3 detected was due to condensation from vapour phase during cooling from the final temperature. The oxidation in the bottom-up furnace at 1600 °C and the arc jet in air resulted in a very similar layered morphology, despite the different temperature and oxygen pressure. On the surface, the dominant phase is the platelet-shaped complex oxide with stoichiometry Hf6Ta2O17 embedded in a borosilicate phase which contains Ta and Hf (see EDS spectrum A in Fig. 3c), under which a porous layer containing Ta-doped hafnia grains is generated. The mechanism of formation of Hf6Ta2O17 phase is not clear. Direct formation of this phase by oxidation of the solid solution is not possible as the content of Ta is too low to reach the exact stoichiometry of the complex oxide. Thus we must conclude that further Ta incorporation in hafnia grains occurred. On the basis of the microstructural observation, it is argued that at P 1600 °C a borosilicate liquid phase containing some dissolved Hf and Ta exists on the surface. Progressively, whilst reactions (1)-(3) occur, the surface borosilicate liquid layer undergoes Ta enrichment due to oxidation of TaSi2. Large convective fluxes are responsible for movement of liquids from the subsurface layer up to the surface, where B2O3 evaporation creates bubbles and craters. Since the surface liquid is rich in Ta, the preferential crystal phase which will crystallize, either upon cooling or when the liquid viscosity reaches a solubility limit, is the intermediate oxide Hf6Ta2O17. At the same time, when a continuous Ta-borosilicate layer is formed on the surface, a porous layer is generated underneath. This could form either via capillarity extraction of the silica liquid towards the surface layer or from formation of SiO(g) as indicated in other works [16,21] by active oxidation of TaSi2: TaSi2 þ O2 ðgÞ ¼ 2SiOðgÞ þ Ta  ð5Þ  Indeed, according to thermodynamic calculations, reaction (5) is favourable at 1600 °C when the oxygen partial pressure is lower than 103 Pa. Subsurface oxidation involving upwelling of liquid borosilicate phases has dramatic effects at temperatures over 2000 °C, as observed for the conical specimen. At these temperatures, ablation/evaporation of the surface liquid can explain the reduced thickness of the external layer. At the same time volume increase due to formation of the Hf6Ta2O17 phase generates large stresses, resulting in scale spallation. It is worth comparing the oxidation behaviour of HfB2-MoSi2 materials with the composite used in the present work. MoSi2-doped HfB2 composites show good stability up to 1900 °C [8,9], with formation of a compact scale with no porosity and adhering well to the bulk. TaSi2 offers good protection up to 1500-1600 °C [8]. This is further proved by the fact that the strength after oxidation is  not significantly varied compared to the pristine value. However, at higher temperatures important issues are not addressed with the addition of this silicide. One concern is the instability of HfO2 scale, which ultimately results in phase transformation upon cooling and oxide cracking and detachment. Ta was indicated as an ideal candidate to modify the HfO2 scale, because it can stuff oxygen into the lattice and induce phase stabilization [15]. However, the present experiments have shown that for a 15 vol.% TaSi2-containing composite, the most evident microstructure evolution is the formation of the Hf6Ta2O17 phase which tends to detach from the underneath layer when the temperature exceeds 2000 °C. The second concern is the melting temperature of this phase, which is not known but could be lower than that of pure HfO2. Finally, TaSi2, unlike MoSi2 [9], seems to promote the formation of a subsurface porous layer which further weakens the scale.  The support of R. Savino and A. Di Maso (DIAS, University of Naples) for arc-jet tests is gratefully acknowledged.  J. Eur. Ceram. Soc.  [1] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, J. Eur. Ceram. Soc. 19 (1999) 2405. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, J. Am. Ceram. Soc. 90 (2007) 1347. [3] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, J. Mat. Sci. 39 (2004) 5925. [4] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, J. Mat. Sci. 39 (2004) 5925. [5] F. Monteverde, A. Bellosi, J. Eur. Ceram. Soc. 25 (2005) 1025. [6] R. Klein, M. Desmaison-Brut, J. Desmaison, L. Mazerolles, M. F. Trichet, High Temperature Corrosion and Protection of Materials 6, Parts 1 and 2, vol. 461-464, 2004, pp. 849-856. [7] D. Sciti, L. Silvestroni, A. Bellosi, J. Mat. Res. 21 (2006) 1460. [8] D. Sciti, A. Balbo, A. Bellosi, (2009) 1809. [9] R. Savino, M.De Stefano Fumo, L. Silvestroni, D. Sciti, J. Eur. Ceram. Soc. 28 (2008) 1899. [10] J.B. Berkowitz-Mattuck, J. Electrochem. Soc. 113 (1966) 908. [11] T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, Acta Mater. 55 (2007) 5999. [12] D. Sciti, L. Silvestroni, G. Celotti, C. Melandri, Guicciardi, J. Am. Ceram. Soc. 91 (2008) 3285. [13] E.J. Opila, S. Levine, J. Lorincz, J. Mater. Sci. 39 (2004) 5969. [14] I.G. Talmy, J.A. Zaykoski, M.M. Opeka, A.H. Smith, J. Mater. Res. 21 (2006) 2593. [15] S. R. Levine, E. J. Opila, NASA/TM-2003-212483. [16] F. Peng, R.F. Speyer, J. Am. Ceram. Soc. 91 (2008) 1498. [17] F. Peng, Y. Berta, R.F. Speyer, J. Mater. Res. 24 (2009) 1855. [18] R.S. Roth, J.L. Wang, Phase Diagram for Ceramists 4392 (1970). [19] C. Zhao, T. Witters, P. Breimer, J. Maes, M. Caymax, S. De Gendt, Microelectr. Eng. 84 (2007) 7. [20] S.N. Karlsdottir, J.W. Halloran, C.E. Henderson, J. Am. Ceram. Soc. 90 (2007) 2863. [21] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Ceram. Soc. 89 (2006) 3240.  J. Am.  29  S.  \\x0c']"
},{
  "_id": 160,
  "PDF": "Oxidation behaviour of zirconium diboride–silicon carbide ceramic composites under low oxygen partial pressure.pdf",
  "Text": "['Corrosion Science 53 (2011) 3742-3746  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Oxidation behaviour of zirconium diboride-silicon carbide ceramic composites under low oxygen partial pressure  Chengyi Tian a, Dong Gao a, Yue Zhang a,⇑ , Chunlai Xu b, Yang Song b, Xiaobin Shi b  a Key Laboratory of Aerospace Materials and Performance (Ministry of Education), School of Materials Science and Engineering, Beihang University, Beijing 100191, PR China b National Key Laboratory of Advanced Functional Composite Materials Technology, Beijing 100076, PR China  a r t i c l e  i n f o  a b s t r a c t  The oxidation behaviour of ZrB2-based ceramics under low oxygen partial pressure ranged from 0.5 to 1.5 kPa was investigated. Low oxygen partial pressure was found to have remarkable effect on phase composition of the surface and the structures of oxide scale. And the thickness and microstructures of oxide scale was characterized by using scanning electron microscopy and X-ray diffraction analysis. The results indicate that the oxidation mechanism of ZrB2-based ceramics changes under low oxygen partial pressure, and the oxidation resistance increases with the reduction of oxygen partial pressure. Ó 2011 Elsevier Ltd. All rights reserved.  Article history:  Received 25 January 2011 Accepted 9 July 2011 Available online 20 July 2011  Keywords:  A. Ceramic matrix composites C. Oxidation B. SEM  1. Introduction  Ultra-high temperature ceramics (UHTCs) recently have been of great interest, motivated by the urgently needs of reusable Thermal Protective System (TPS) and other components for future generation of hypersonic aerospace vehicles [1]. Of those UHTCs, ZrB2 is widely investigated due to its unique properties such as high melting point (>3000 °C), low density (6.09 g/cm3), high strength, and high-elastic modulus [2,3]. In addition, it was found that the introduction of second phase particles significantly improved the mechanical properties and oxidation resistance performance. Thus ZrB2-based ceramics are considered to be the potential candidate materials used in TPS for re-entry aircrafts and supersonics vehicles [4-7]. Oxidation resistance performance is most challenging in the development of ZrB2-based ceramics, because it directly determines whether the materials could survive under oxidizing condition at high temperatures. Great efforts have been devoted to improve the oxidation resistance performance of ZrB2-based materials under high temperature, and the most common strategy is to introduce the additives such as SiC, MoSi2, or TaSi2 [8-10]. It has been shown that the introduction of SiC into the ZrB2 ceramics leads to significant improvement in oxidation resistance performance by encouraging the formation of a borosilicate glass layer on exposed surfaces. And ZrB2-SiC ceramic composites are currently considered as the baseline UHTCs [11-14]. Most of the investigations of the oxidation behaviour of ZrB2-SiC ceramics were conducted in either static or flowing dynamic air un ⇑ Corresponding author. Tel./fax: +86 10 82316976. E-mail address: zhangy@buaa.edu.cn (Y. Zhang).  0010-938X/$ see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2011.07.020  der ambient pressures. In order to understand the oxidation progress of the materials in re-entry environments, it is in urgent need to investigate the behaviour of the ceramic composites under low oxygen partial pressure. Recently the oxidation resistance properties of the ceramics in simulated re-entry condition were studied [15]. And some novel methods of testing UHTCs at high temperature were used, such as arc jet method and ribbon method [16,22]. Han et al. [17], Fahrenholtz et al. [18] and Gao et al. [19], respectively reported the oxidation behaviour of ZrB2-SiC under low oxygen partial pressure. Han investigated oxidation resistance performance of ZrB2-SiC under air pressure with oxygen partial pressure of 2 \\x02 10\\x004 atm. And the results indicated lower oxygen partial pressure was detrimental for the oxidation resistance property of the ceramics in comparison with those ceramics oxidized in air. Fahrenholtz focused on the effect of oxygen partial pressure on the evolution of the oxidation layer and SiC layer, and proposed a detailed explanation under reducing condition produced by CO and CO2 as well. Gao’s work premised on the formation mechanism of zircon during oxidation of ZrB2-SiC ceramics under a low oxygen partial pressure of 0.2 kPa. According to his research, lower oxygen partial pressure was beneficial to the formation of zircon due to the active oxidation of SiC. However, compared with the oxidation of ceramics at ambient pressures, the oxidation mechanism at low oxygen partial pressure was rarely investigated. Therefore, to attain a more comprehensive and systematic understanding of both oxidation and evolution of oxidation products under various low oxygen partial pressures is desirable. The purpose of this paper is to describe the oxidation behaviour of ZrB2-SiC under low oxygen partial pressures. The composition of surface layer and variation thickness of oxide scale were investigated to analyse oxidation resistance of ceramic composites.  \\x0c', 'C. Tian et al. / Corrosion Science 53 (2011) 3742-3746  3743  2. Experimental procedure  2.1. Materials processing  3. Results and discussion  3.1. Surface microstructure evolution  Commercially available ZrB2 and SiC powders with a reported purity 99% and an average particle size of 2 lm were used to prepare the materials for this study. ZrB2-based ceramics containing of 20 vol% SiC (ZS2) were prepared by hot pressing method, which is referred as ZS2 in the following discussion. The mixtures were hotpressed at 1900 °C for 1 h at the pressure of 30 MPa, in a graphite die lined with graphite foil coated with BN. After loading to graphite dies, the furnace was heated to targeted temperature at an average rate of 30 °C/min by using argon as protective gas. The load was removed when the die temperature dropped below 1750 °C. After hot-pressing time elapsed, the furnace was cooled to the room temperature. The sintered specimens were taken out from the graphite dies. Then the specimens with dimensions of 10 \\x02 10 \\x02 6 mm were cut from sintered materials by electrical discharged machining (EDM) method. And the surface of the specimens was diamond polished to 1 lm finish. Before oxidation test, the coupons were ultrasonically cleaned successively in deionized water and ethanol.  2.2. Oxidation testing  A tube furnace with resistive molybdenum disilicide (MoSi2) heating element up to 1800 °C was used for the oxidation studies. Oxygen (>99.99%) and Nitrogen (>99.99%) were introduced at the end of the alumina tube. The total pressure was fixed at about 15 kPa (standard deviation: 0.1 kPa), and the expected oxygen partial pressure can be achieved by controlling the flow rate ration between oxygen and nitrogen. A digital pressure gauge was used to measure and monitor gas pressures in the alumina tube. Specimens were sent to the center of the furnace, placed on an alumina container. The specimens were heated to a certain temperature of 1500 °C at a ramp rate of approximately 5 °C/min by using argon as protective gas. When reaching target temperature, oxygen and nitrogen, which had been adjusted according to the expected ratio (as shown in Table 1), were introduced into the furnace. After the expected oxidation time was achieved, the introduction of oxygen and nitrogen was stopped. And the samples were cooled in nitrogen atmosphere.  2.3. Characterization of oxidized sample  Phase compositions of oxidized specimens were examined by X-ray diffraction (XRD, Dmas-2200, Rigaku, Tokyo, Japan) with Cu Ka radiation. Specimens were prepared for microscopy by cutting cross-section and polished using diamond abrasives before SEM observation. In addition, the microstructure and chemical composition of oxidized specimens on the surface and crosssection oxidation layer were detected by SEM (CamScan-3400, Cambridge, England), which was equipped with an EDS detector (INCAINCAPentaFET-x3, Oxford, England).  Table 1 The oxidation condition of ceramics samples.  Specimens  ZS2-10  ZS2-15  ZS2-20  ZS2-25  ZS2-30  V N2 =V O2 PO2  (kPa) Temperature (°C) Time (min)  9  1.5  1500 30  14  1  1500 30  19  0.75  1500 30  24  0.6  1500 30  29  0.5  1500 30  Fig. 1 shows the X-ray diffraction patterns of ZS2 ceramics after oxidized at 1500 °C for 30 min with various low oxygen partial pressures. From Fig. 1, it can be seen that only ZrB2 and ZrO2 were observed in the oxidized samples (ZS2-20, ZS2-25 ZS2-30), suggesting the three samples were slightly oxidized. ZrB2 phase could be detected due to thin oxide scale, as shown in Fig. 2. Since the three ceramic samples share the same composition and the main product on the surface is zirconia, the size and amount of zirconia should represent the oxidation degree of the ceramic composites. As shown in the image, the relative intensity of the zirconia decreases with reduction of the partial oxygen pressure from 0.75 to 0.5 kPa. This indicates the oxidation resistance of ceramics is improved under reducing partial oxygen pressure. Compared with the above three samples, only ZrO2 could be identified from oxidized ZS2-15 and the ZrB2 peaks disappeared (Fig. 1 ZS2-15), which means the oxidation resistance of ZS2-15 ceramics decays. XRD analysis revealed that zircon peak could be detected besides zirconia (Fig. 1 ZS2-10). According to formation mechanism of zircon [19], the probable formation routes of zircon are shown in Eqs. (1) and (2). And the latter is related with the active oxidation of SiC, expressed as Eq. (3). The relative intensity of the zirconia of ZS2-10 was almost the same as that of ZS2-15; meanwhile there was a lot of zircon formed on the ceramics surface of ZS2-10. This suggests more zirconia is formed on the surface of ZS2-10 and the oxidation resistance of ZS2-10 is further deteriorated in comparison to ZS2-15. With the reduction of oxygen partial pressure, the experimental results indicate the amount of formed zirconia decreases, but the oxidation resistance of ZS2 ceramics increases.  ZrO2 ðcrÞ þ SiO2 ðlÞ ! ZrSiO4 ðcrÞ  ZrO2 ðcrÞ þ SiOðgÞ þ 1=2O2 ðgÞ ! ZrSiO4 ðcrÞ  SiCðcrÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  3.2. Cross-section microstructure evolution  ð1Þ  ð2Þ  ð3Þ  Fig. 2 shows the cross-section SEM image of oxidized materials under various low oxygen partial pressures. As shown in the image,  Fig. 1. X-ray diffraction of ZrB2-20 vol% SiC ceramics after oxidized at 1500 °C for 30 min under various oxygen partial pressure.  \\x0c', '3744  C. Tian et al. / Corrosion Science 53 (2011) 3742-3746  Fig. 2. Backscatter electron images of the cross section of ZrB2-20 vol% SiC ceramics oxidized at 1500 °C for 30 min under various oxygen partial pressures (a) ZS2-10 (PO2 = 1.5 kPa); (b) ZS2-15 (PO2 = 1 kPa); (c) ZS2-20 (PO2 = 0.75 kPa); (d) ZS2-25 (PO2 = 0.6 kPa) and (e) ZS2-30 (PO2 = 0.5 kPa).  the oxidized specimens (including ZS2-15, ZS2-20, ZS2-25, and ZS2-30) produced a two layered structure: (a) a rich-in zirconia, containing little silicate; (b) unaffected ZrB2-SiC. The two layered structure is consistent with the report by Rezaie et al. [18]. In our research, the ceramic composites were oxidized under oxygen pressure between 0.5 and 1.5 kPa. In this case, it should lead to both active oxidation and passive oxidation of SiC according to the volatility diagram ZrB2-SiC system [18]. As a result, both the vapor pressures of SiO and SiO2 were relatively high. In this case, it is not easy to form the protective oxide layer on the ceramics surface. It should be noted that the total pressure was 15 kPa, which furthermore prevented the formation of silicate layer. Also, the ZrB2 would be oxidized to form ZrO2 and B2O3. The latter was volatile, thus only ZrO2 remained as the stable condensed phase. Because there was no protective layer on the surface of oxidized specimen, oxygen partial pressure beneath zirconia layer was too high to form the SiC-depleted layer. As a result, the two distinct layers were formed. As for ZS2-10, four layers could be identified according to Fig. 2, composed of a silica-rich glass, a zriconia-rich oxidized layer, a SiCdepleted layer (note that SiC was only partially depleted and not fully removed) and unaffected ZrB2-SiC. The structure with four layers is the same as that of samples oxidized at air pressure, observed by other investigators [17,23,24]. The formation of SiC-depleted layer of ZS2-10 was mainly attributed to active oxidation of SiC beneath the zirconia-rich layer. Zirconia-rich layer was formed due to oxidation of ZrB2 and passive oxidation of SiC. Because of the active oxidation of SiC beneath zirconia-rich layer and on the surface of specimen, the gas partial pressure (SiO and SiO2) was gradually increasing. As a result, silicate glass was found on the surface of the sample. In addition, the gaseous SiO reacted with O2 to form SiO2 according to Eq. (4), which further promoted the formation of silicate glass on the surface. As a result the four layered structure can be observed in ZS2-10 sample.  2SiOðgÞ þ O2 ðgÞ ! 2SiO2 ðlÞ  ð4Þ  The structure of specimen oxidized under oxygen pressure of 1.5 kPa was similar to that oxidized in air, while the structures of other four specimens were just like those oxidized under extremely low oxygen partial pressure. It means the oxidation mechanism  started to change during oxygen partial pressure between 1 and 1.5 kPa. When the oxidation of ZrB2-SiC produced the structure consisted of four layers, the oxidation process is consistent with parabolic mass gain kinetics based on the presence of a protective layer reported by Fahrenholtz [21]. Otherwise, the oxidation process is consistent with linear mass gain kinetics and reaction ratecontrolled kinetics [18], when oxidation of ZrB2-SiC produced two distinct layers. Therefore, the oxidation mechanism of ZS2 ceramics transforms from parabolic mass gain kinetics to linear mass gain kinetics. The variation of thickness of the oxide scale under different oxygen partial pressure could be identified from Fig. 3. The thickness of the oxidized layer for ZS2 ceramics decreased with the reduction of oxygen partial pressure, suggesting that low oxygen partial pressure is beneficial to oxidation resistance property of samples in this oxygen range. According to the volatility diagram of ZrB2-SiC system, the low oxygen partial pressure is beneficial to the active oxidation of SiC. The two layered structure can be observed because of the active oxidation of SiC. The stability of  Fig. 3. Thickness of the oxide scale of ZrB2-20 vol% SiC ceramics oxidized at 1500 °C for 30 min under various oxygen partial pressures.  \\x0c', 'C. Tian et al. / Corrosion Science 53 (2011) 3742-3746  3745  Fig. 4. Backscatter electron images of the surface of ZrB2-20 vol% SiC ceramics oxidized at 1500 °C (a-c) ZS2 oxidized under oxygen partial pressures of 1 kPa for various times min (a) 20 min; (b) 60 min; (c) 90 min; (d-f) ZS2 oxidized under oxygen partial pressures of 1.5 kPa for various times (d) 20 min; (e) 60 min; (f) 90 min.  the oxide scale is due to formation of a ‘‘solid pillars, liquid roof’’ [20]. And the solid phase (ZrO2) and silica-rich liquid phase is tightly held by capillary forces. In fact, ZrO2 phase is often highly porous and silica-rich liquid usually is not uniformly dispersed, so it is difficult to combine tightly for solid and liquid phase. It means pores could inevitably exist in the oxide scale. We infer that fewer pores are formed with the reduction of oxygen partial pressure in the oxide scale. Fewer pores mean oxidation rate of specimens is slower. Thus the thickness of the oxidized layer for ZS2 ceramics is decreasing with the reduction of oxygen partial pressure.  3.3. Oxidation kinetics behaviour  To further verify the analysis based on the oxide scale thickness as a function of oxygen partial pressure, the evolution of oxide scale was also investigated at different time under various oxygen partial pressures. When ZS2 specimens were oxidized under oxygen partial pressure of 1 kPa, there were two different kinds of oxide scale structure (Fig. 4(a-c)). One was composed of two layers: (a) a rich-in zirconia, containing little silicate; (b) unaffected ZrB2-SiC; while the other contained three layers: a zirconia-rich oxidized layer, a SiC-depleted layer and unaffected ZrB2-SiC.  ZrB2 ðcrÞ þ 5=2O2 ðgÞ ! ZrO2 ðcrÞ þ B2O3 ðlÞ  SiCðcrÞ þ 3=2O2 ðgÞ ! SiO2 ðgÞ þ COðgÞ  ð5Þ  ð6Þ  From the structure of ZS2 specimens, we can infer the evolution of oxide scale with time. At the beginning of oxidation period, the ZrB2 and SiC were respectively oxidized by Eqs. (5) and (6). With the increasing of oxidation time, the surface of specimens was covered by dense layer of zirconia particles, which silica glass distributed in the grain boundary. Thus the zirconia-rich layer was formed (Fig. 4(a)). However, the amount of silica glass distributed in the zirconia layer began to grow with the extending of oxidation time. Thus the oxygen partial pressure in the inner part of the ceramics was lowered. As a result, the active reaction of SiC began and the SiC-depleted layer was formed. It was easy for silica glass to evaporate at elevated temperature of 1500 °C, so there were several large black holes in zirconia-rich layer (Fig. 4(c)). Furthermore,  the active reaction of SiC could produce amounts of CO gas (Eq. (6)). They could gather in the zirconia-rich layer and sharply increase the formation of hole as well. ZS2 specimens oxidized under oxygen partial pressure of 1.5 kPa, there were also two different kind of oxide scale structure (Fig. 4(d- f)). One was that consisted of two layers: (a) a rich-in zirconia, (b) unaffected ZrB2-SiC; the other one consisted of four layers: a silica-rich layer, a zirconia-rich oxidized layer, a SiC-depleted layer and unaffected ZrB2-SiC. We can infer the evolution of oxide scale by the structure of ZS2 specimens at different time. At the first period of oxidation, Zirconia-rich layer was formed due to the oxidation of ZrB2 and passive oxidation of SiC. There was no silica glass formed, which must be likely due to the volatilization of SiO and SiO2 under low atmospheric pressure. With gas (SiO and SiO2) partial pressure increasing, the silica-rich layer firstly appeared at 30 min. The silica-rich glass layer became thicker with the oxidation of specimens and increasing of time. Since the silica-rich glass layer formed, the oxygen partial pressure beneath zirconia-rich layer was so low that the active reaction of SiC occured. Amounts of black holes were  Fig. 5. Variation of thickness of the oxide scale for ZrB2-20 vol% SiC ceramics oxidized at 1500 °C under oxygen partial pressures of 1 kPa.  \\x0c', '3746  C. Tian et al. / Corrosion Science 53 (2011) 3742-3746  rate-controlled kinetics. Our experimental results verified that the oxidation mechanism of ZrB2-SiC ceramics changed during oxygen partial pressure between 1 and 1.5 kPa. In addition, the results of X-ray diffractions, and cross-section images of samples confirmed that the oxidation resistance of ZrB2-SiC ceramics increased with the reduction of oxygen partial pressure under low oxygen partial pressure.  Acknowledgements  This work is financially supported by the Cheung Kong Scholars and Innovative Research Team Program in University from Ministry of Education under Grant No. IRT0805.  References  Fig. 6. Variation of thickness of the oxide scale for ZrB2-20 vol% SiC ceramics oxidized at 1500 °C under oxygen partial pressures of 1.5 kPa.  found in the zirconia layer because of evaporation of silica glass (Fig. 4(f)). Comparing the structure evolution of ZS2 under oxygen partial pressure of 1 and 1.5 kPa, we could find the specimen oxidized under oxygen partial pressure of 1 kPa did not form silica-rich layer for a long time. However, silica-rich layer was found to form in the specimen oxidized under oxygen partial pressure of 1.5 kPa for 30 min or longer time. The variation of the thickness of oxide scale as a function of holding time at 1500 °C under oxygen pressure of 1 and 1.5 kPa are respectively shown in Figs. 5 and 6. And they can be an indicator of the oxidation mechanism. When ZS2 specimens oxidized under oxygen partial pressure of 1 kPa, the linear trend indicates the reaction rate-controlled kinetics. While, that under oxygen partial pressure of 1.5 kPa, the parabolic trend suggests that the oxidation of ceramic composites is diffusion rate-controlled kinetics. It means the oxidation mechanism of ZS2 ceramics begins to change during oxygen partial pressure between 1 and 1.5 kPa. This is consistent with previous analysis. Because silica glass layer did not form on the surface of ceramic specimens, the thickness of oxide scale sharply increased with a linear trend as shown in Fig. 5. However, owing to the formation of silica-rich glass layer, the thickness of oxide scale increased with a parabolic trend (Fig. 6). At the beginning of the oxidation, the thickness of oxide scale sharply increased, and the trend was almost linear. This is because no silica-rich glass layer is formed on the surface of the sample (Fig. 4(d)). With the formation of silicarich layer, the oxidation rate started to gradually decrease.  4. Conclusion  Specimens of ZrB2-based ceramics containing 20 vol% SiC were oxidized under low oxygen partial pressure ranged from 0.5 to 1.5 kPa. When ZrB2-SiC ceramic specimens were oxidized under oxygen partial pressure of 1.5 kPa, it produced a structure consisted of four layers, and the variation of the thickness of oxide scale as a function of holding time was in parabolic trend. The results suggest that the oxidation of ceramic composite is the diffusion rate-controlled kinetics. While the ceramic specimens were oxidized under oxygen partial pressure of equal or lesser than 1 kPa, it produced two distinct layers, and the variation of the thickness of oxide scale as a function of holding time was in linear trend. It indicates the oxidation process is consistent with reaction  [12]  [1] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000 °C + hypersonic aerosurface. Theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887-5904. [2] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. 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Davis, H.C. Graham, Effect of an SiC addition on the oxidation of ZrB2, Am. Ceram. Soc. Bull. 52 (1973) 612-616. [9] D. Sciti, M. Brach, A. Bellosi, Long-term oxidation behavior and mechanical strength degradation of a pressurelessly sintered ZrB2-MoSi2 ceramic, Scripta Mater. 53 (2005) 297-1302. [10] F. Peng, R.F. Speyer, Oxidation resistance of fully dense ZrB2 with SiC, TaB2, and TaSi2 additives, J. Am. Ceram. Soc. 91 (2008) 1489-1494. [11] S.S. Hwang, A.L. Vasiliev, N.P. Padture, Improved processing and oxidationresistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids, Mater. Sci. Eng. A 464 (2007) 216-224. I.B. Ban’kovskaya, V.A. Zhabrev, Kinetic analysis of heat resistance of ZrB2-SiC composites, Glass Phys. Chem. 31 (2005) 482-488. [13] W.B. Han, P. Hu, X.H. Zhang, J.C. Han, S.H. Meng, High-temperature oxidation at 1900 °C of ZrB2-xSiC ultrahigh-temperature ceramic composites, J. Am. Ceram. Soc. 91 (2008) 3328-3334. [14] D. Sciti, M. Brach, A. Bellosi, Oxidation behaviour of a pressureless sintered ZrB2-MoSi2 ceramic composites, J. Mater. Res. 20 (2005) 922-930. [15] F. Monteverde, R. Savino, Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmosphere re-entry conditions, J. Eur. Ceram. Soc. 27 (2007) 4797-4805. J.C. Han, P. Hu, X.H. Zhang, S.H. Meng, W.B. Han, Oxidation-resistant ZrB2-SiC composites at 2200 °C, Compos. Sci. Technol. 68 (2008) 799-806. J.C. Han, P. Hu, X.H. Zhang, S.H. Meng, Oxidation behavior of zirconium diboride-silicon carbide at 1800 °C, Scripta Mater. 57 (2007) 825-828. [18] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Oxidation of zirconium diboride- silicon carbide at 1500 °C at a low partial pressure of oxygen, J. Am. Ceram. Soc. 89 (2006) 3240-3245. [19] D. Gao, Y. Zhang, J.Y. Fu, C.L. Xu, Y. Song, X.B. Shi, Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: formation of zircon, Corros. Sci. 52 (2010) 3297-3303. J. Li, T.J. Lenosky, Thermochemical and mechanical stabilities of the oxide scale of ZrB2 + SiC and oxygen transport mechanisms, J. Am. Ceram. Soc. 91 (2008) 1475-1480. [21] W.G. Fahrenholtz, Thermodynamic analysis of ZrB2-SiC oxidation: of a SiC-depleted region, J. Am. Ceram. Soc. 90 (2007) 143-148. [22] S.N. Karlsdottir, J.W. Halloran, Rapid oxidation characterization of ultra-high temperature ceramics, J. Am. Ceram. Soc. 90 (2007) 3233-3238. [23] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics, J. Mater. Sci. 39 (2004) 5925-5937. [24] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500 °C, J. Eur. Ceram. Soc. 27 (2007) 2495-2501.  formation  [16]  [17]  [20]  \\x0c']"
},{
  "_id": 161,
  "PDF": "Oxidation behaviour of ZrB2–SiC (Al-Y) ceramics at 1700 °C.pdf",
  "Text": "['Journal of the European Ceramic Society 36 (2016) 3769-3774  Contents lists available at www.sciencedirect.com  Journal  of  the  European  Ceramic  Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Oxidation behaviour of ZrB2-SiC (Al/Y) ceramics at 1700  C  Jiabei He a , Yiguang Wang a,∗ , Lei Luo a , Linan An b,∗  a School of Materials Science and Engineering, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China b Advanced Materials Processing and Analysis Center, University of Central Florida, Orlando, FL 32816, USA  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  ZrB2 -SiC(Al/Y) ceramic composites were prepared by hot-pressing a mixture of ZrB2 and polymer-derived SiC doped with either aluminium or yttrium (SiC(Al/Y)). The obtained composites were oxidized at 1700   C in air. The weight change as a function of the oxidation time was measured to study the oxidation kinetics. It was found that the ZrB2 -SiC(Al/Y) composites exhibited significantly improved oxidation resistance, as compared to the composite of ZrB2 with un-doped SiC. The structure of the oxide resulted from the oxidation of the composites was characterized by scanning electron microscopy and X-ray diffraction. The results suggested that the improvement can be attributed to the formation of yttrium/aluminium-doped silica and/or yttrium disilicate, which effectively protect the composites. © 2016 Elsevier Ltd. All rights reserved.  Article history:  Received 28 December 2015 Received in revised form 19 February 2016 Accepted 19 February 2016 Available online 2 March 2016  Keywords:  ZrB2 -SiC Doped Oxidation kinetics Oxidation Oxidation resistance  1.   Introduction  In recent years, ultra-high temperature ceramic (UHTC) materials have played an  important role  in the  fabrication of thermal protection systems (TPSs)  for hypersonic aerospace vehicles due to  their good adaptability  for critical applications  [1-3]. UHTCs, including refractory metal diborides and carbides, have been studied since  the 1960s  [4-6]. These ceramics can be used  in harsh environments, for instance, during atmospheric reentry and hypersonic flights  [2,7], due to their high melting temperature (above 3000  C), good ablation resistance, and high mechanical strength at high temperatures. Of all UHTC materials, ZrB2 has received the special attention because it offers the lowest density (6.09 g/cm3 ) in addition to a high melting point and a good thermal shock resistance  [8,9]. However,  its high-temperature properties are  likely to deteriorate in oxygen-rich environments. Therefore, the oxidation behaviour of ZrB2 and ZrB2 -based ceramics is of great concern [8,10,11]. The oxidation of ZrB2 in air usually  leads to the  formation of ZrO2 and  liquid B2O3 [12,13]. However, because of  the evaporation of boron oxide at  temperatures  in  the  range  from 1100  to 1400  C, ZrB2 is usually combined with other  refractory ceramics  like SiC  for high-temperature applications  [14-17]. ZrB2 -SiC exhibits a good oxidation resistance at temperatures  lower than  ∗ Corresponding authors. Fax: +1 4078230208. E-mail addresses: wangyiguang@nwpu.edu.cn (Y. Wang), linan.an@ucf.edu (L. An).  http://dx.doi.org/10.1016/j.jeurceramsoc.2016.02.037 0955-2219/© 2016 Elsevier Ltd. All rights reserved.  1700  C due to the formation of a triple-layered structure consisting of a borosilicate/silica outer-layer, a ZrO2 -SiO2 middle-layer, and a SiC-depleted  inner-layer [11]. As the flight speed of hypersonic vehicles further and further increases, the TPS will eventually face an oxidative environment, with temperatures exceeding 1700  C [18]. The practical application of SiO2 -forming ceramics has been temperature of 1725  C reported  to be  limited  to a maximum  due to the occurrence of a rapid oxidation, as well as the potential volatility, scale melting, and scale/substrate reactions of SiO2 [19]. Therefore, the borosilicate/silica outer-layer of ZrB2 -SiC may quickly evaporate, resulting in the formation of a porous ZrO2 layer and fast failure of the TPS [2,7,19]. Finding methods to improve the in air at temperatures higher than 1700  C  stability of ZrB2 -SiC  is therefore of great scientific and practical  interest. However, very few studies  investigating this phenomenon have been conducted and reported so far. The maximum application temperature of ZrB2 -SiC can either be improved by reducing the silica activity at high temperatures or forming silicates on the top of the ceramic composites. In a previous study, we tried to dope SiC with aluminium in order to reduce the silica activity at high temperatures [12], and it was found that the oxidation resistance of silicon-based ceramics was enhanced at temperatures of 1200  C [20,21]. Even at temperatures of up to 1500  C, the silica activity was observed to be reduced, although the  reduction was more prominent at  the  lower  temperatures [12,22].  It  is also well known  that  the out-diffusion of elements like yttrium during the oxidation of sintered Si3N4 , which results in the formation of yttrium-containing silica, can  improve the hightemperature oxidation resistance of Si3N4 [23-25]. Therefore,  in                                              \\x0c', '3770   J. He et al. / Journal of the European Ceramic Society 36 (2016) 3769-3774  Fig. 1. Molecular structures of PCS (a); PACS (b); and PYCS (c). These structures are provided by the manufacturer.  this study, we adopted this method and doped SiC with aluminium or yttrium to form SiC(Al/Y), then introduced the SiC(Al/Y) into the ZrB2 -SiC compounds in order to improve the oxidation resistance of ZrB2 -based ceramics at 1700  C.  2. Material and methods  Commercially  available  ZrB2 (0.5  \\u242em,  99%  purity,  Beijing Mountain Technical Development Center, Beijing, China),  solid polycarbosilane  (PCS, Laboratory of Special Advanced Materials, Xiamen University, Xiamen, China), solid polycarbosilane containing 2 or 7 wt% yttrium (PYCS-2 and PYCS-7, respectively, Laboratory of Special Advanced Materials, Xiamen University, Xiamen, China), and solid polycarbosilane containing 2 wt% aluminium (PACS, Laboratory of Special Advanced Materials, Xiamen University, Xiamen,  Fig. 2. XRD patterns of the as-prepared ZrB2 -SiC (Al/Y) ceramic composites.  China) were used  as  raw materials  for  the  fabrication of  the ZrB2 -SiC and ZrB2 -SiC(Y/Al) ceramic composites. According to the manufacturer, the possible molecular structures of PCS, PACS, and PYCS are shown in Fig. 1. The residual percentages for PCS, PYCS, and PACS were all around 50 wt%, with no obvious difference after the pyrolysis at 1800  C. The chemical composites for PCS, PACS, PYCS2, and PYCS-7 were measured to be SiC1.29O0.05 , SiAl0.02C1.28O0.03 , SiY0.01C1.21O0.02 , and SiY0.04C1.21O0.04 , respectively. The density of the pyrolysed products was considered to be 3.2 g/cm3 , similar to the density of pure SiC. The raw materials were weighed accordingly  to prepare  the ZrB2 -20 vol% SiC and ZrB2 -20 vol% SiC(Y/Al) composites. The SiC and SiC(Y/Al) powders were prepared by adopting the method proposed by  Isihara et al. [26]. The solid precursor (PCS, PYCS or PACS) was first dissolved in hexane to form a 20 wt% solution. The solution was then gradually poured into ethanol to create followed by a curing step at 250  C  a precipitate,  for 10 h. Then, the resulting sample was pyrolysed at 1000  C for 4 h. Fine spherical SiC and SiC(Y/Al) powders were obtained after the pyrolysis. Next,  these powders were mixed with  the ZrB2 powder according  to  the design  ratio,  followed by ball-milling  in an agate  jar with agate balls for 5 h. The resulting mixed powders were poured  Fig. 3. Comparison of the morphologies of the polished surfaces of the different ZrB2 -SiC (Al/Y) ceramic samples: (a) ZS; (b) ZSA; (c) ZSY-2; (d) ZSY-7.  \\x0c', 'J. He et al. / Journal of the European Ceramic Society 36 (2016) 3769-3774   3771  zirconia tube furnace (GSL-1800X, Hefei Kejing Materials Technology Co., Ltd., Hefei, China).  Inside the  furnace, the samples were placed  in a zirconia crucible on  top of a zirconia plate mat and heated to 1700  C at a rate of 10  C/min under a protective highpurity argon atmosphere. After the oxidation temperature had been reached, the samples were stabilized at 1700  C  for several minutes, followed by exposure to pure air at a flow rate of 10 ml/min. After the designated oxidation time (10, 20, 30, 40, 50 and 60 min) had been reached, the samples were cooled down to room temperature under argon atmosphere. The weight gain was determined as a function of the oxidation time using a balance with an accuracy of 0.1 mg (Mettler Tolendo AL204, Greifensee, Switzerland). The experiments were repeated at least three times with different samples for each designated oxidation time. The phase composition of the oxidized samples was analysed by X-ray diffraction (XRD, Rigaku D/max-2400, Tokyo, Japan) with CuK␣  radiation, with the data digitally recorded in continuous scanning mode in the 2\\x02  range from 10 to 80  at a scanning rate of 0.02  s−1 . The morphologies of the as-received samples and the oxidized samples were studied by scanning electron microscopy (SEM, JEOL 6700F, Tokyo,  Japan), and energy-dispersive X-ray spectroscopy (EDS, EDAX, USA) was performed to determine the elemental composition.  3. Results and discussion  Fig. 4.  (a) Comparison of the specific weight change of the different ZrB2 -SiC (Al/Y) ceramics at 1700   C as a function of the oxidation time. (b) The square of the specific weight change as a function of the oxidation time for ZSA, ZSY-2, and ZSY-7.  Fig. 5. XRD patterns of the oxidized ZrB2 -SiC (Al/Y) ceramic composites.  into a graphite die and hot-pressed at 1800  C for 1 h  in an argon atmosphere under a uniaxial pressure of 35 MPa. The obtained ceramic samples prepared from the PCS, PACS, PYCS-2 or PYCS-7 precursor were denoted as ZS, ZSA, ZSY-2 and ZSY-7, respectively. For the oxidation studies, all specimens were cut into uniform pieces with dimensions of 10 mm   5 mm   5 mm. Each surface of the samples was polished to a 0.5 finish using diamond paste. An isothermal oxidation study was then conducted at 1700  C using a  ×  ×  By applying the Archimedes method, the relative densities of the four as-received ceramics were determined to be 95.75% (ZS), 96.66% (ZSA), 96.48% (ZSY-2), 95.93% (ZSY-7), respectively. Fig. 2 shows the XRD patterns obtained for the four different ceramics, indicating that the main phases in these ceramics are ZrB2 and SiC, with a trace of ZrO2 . Any phases containing aluminium or yttrium could not be detected by XRD either because of  the very  low Al and Y contents or because these elements were dissolved into the matrix to form a solid solution. The morphologies of the polished surfaces of  the different ceramics  (Fig. 3) are also  identical. The dark grains consist of SiC and  the grey ones consist of ZrB2 . No other phases were observed by SEM. All ceramics showed a similar surface morphology. The doping with yttrium or aluminium did not have an obvious effect on the phase composition and morphology of the different ZrB2 -SiC ceramics. The specific weight change of the oxidized samples as a function of the oxidation time is shown in Fig. 4(a). The maximum oxidation time was 60 min for all samples except the ZS due to the spallation of the oxide layer after 40 min of oxidation. For the ZS sample, the weight gain was found to be directly proportional to the oxidation time. For the other three SiC samples doped with Al/Y, the weight gain curves follow a parabolic path, as indicated by Fig. 4b, where the plots of the square of the specific weight gain vs. the oxidation time show a linear relationship. The slopes of the fitting curves in Fig. 4b were considered as the oxidation rate constants, denoted as kp , which were then used to evaluate the oxidation rate of the mate 10−5 , 1.17   10−5 , rials. The kp values were calculated to be 1.58   10−5 g2 /(cm4 × and 0.13   min) for ZSA, ZSY-2 and ZSY-7, respectively. The fitting curve obtained for ZSA may be passing through the origin, whereas an extrapolation of the fitting curves obtained for both ZSY-2 and ZSY-7 suggests an  intercept with the positive part of the y-axis. This means that the initial oxidation of the ZSY2 and ZSY-7 samples proceeds very rapidly,  followed by a  lower oxidation rate later on, indicating that there are two different mechanisms at work during the  initial stage and the  later stage of the oxidation process. For the ZSA samples, the oxidation mechanism remained the same over the entire oxidation period, as  indicated by the uniform oxidation rate.  ×  ×  ×  \\x0c', '3772   J. He et al. / Journal of the European Ceramic Society 36 (2016) 3769-3774  Fig. 6. Comparison of the cross-sectional morphologies of the oxidized ceramics: (a) ZS; (b) ZSA; (c) ZSY-2; (d) ZSY-7.  Table 1 Comparison of the thickness of the oxide scales and the oxidation rate constant kp obtained for the different ZrB2 -SiC ceramic composites after oxidation at 1700   C in pure air for 1 h.  Sample   ZSA  ZSY-2  ZSY-7   Thickness of the oxide scale (\\u242em)   Thickness of the SiC-depleted layer (\\u242em)   ± ± ±   20   20   10   200  500  200   ± ± ±   20   30   15   150  250  100   Oxidation rate constant kp (g2 /(cm4 ×  min))  10−5  10−5  10−5  1.58  1.17  0.13   × × ×  The XRD patterns of the samples after 30 min of oxidation are presented in Fig. 5. No silica peaks were detected when scanning the surface of the ZS samples, whereas the ZSY-2, ZSY-7, and ZSA samples showed strong silica peaks. Compared with the other samples, additional Y2 Si2O7 peaks were found when studying the surface of the ZSY-7 samples. In previous studies [10-15], the formation of a top silica layer was observed on the oxidized ZrB2 -SiC samples at temperatures of up to 1500  C. However, in this study, we did not find silica on the surface of the oxidized ZS samples. It is assumed that the silica quickly evaporates at 1700  C [2], resulting in a linear oxidation behaviour (Fig. 4(a)). The quick loss of silica can also lead to the  formation of a pure zirconia  layer on the top surface after a certain period of time, resulting in spallation once a critical thickness is reached. The oxidation behaviour of ZSA  is different compared with ZS although just a small amount of aluminium has been doped into the SiC. According to previous studies [12,20,21,27,28], the silica oxide formed on the surface contains a small amount of aluminium, which does not interrupt the silica network, but could reduce the activity of silica. At 1700  C, the silica  layer still exists on top of the ZSA samples, preventing  further oxidation.  In this case, the oxidation rate follows a parabolic path, and the oxide  layer remains stable, even after 60 min of oxidation.  In contrast,  for  the ZSY samples,  two different oxidation mechanisms might be at work during the oxidation process. Since Y2 Si2O7 was detected on the ZSY-7 samples, it is believed that there is an out-diffusion of yttrium towards the surface during oxidation. This process was also observed when studying the oxidation of sintered Si3N4 , in which the sinter-aids, e.g., aluminium and yttrium, diffused to the surface during oxidation [29]. Studies on the oxidation of ZrB2 -SiC mixed with other compounds, e.g., TaC, also  indicated that the out-diffusion of Tacontaining species to the surface was the rate-controlling process [10]. Therefore, the diffusion of yttrium to the surface is considered the  initial oxidation mechanism occurring  in the ZSY samples. At the beginning, the accumulation of yttrium at the surface leads to a deterioration of the silica network, resulting  in a fast oxidation during the  initial stage of the oxidation process. Like aluminium, the yttrium may be dissolved  in the silica matrix, thereby reducing  its activity at high temperatures. However, the dissolution of yttrium in silica is limited. When the yttrium concentration at the surface reaches a certain level, the Y2 Si2O7 phase could be formed. The formation of a protective silica layer at the surface lowered the oxidation rate. In this case, the oxidation rate becomes slower compared to the initial stage of the oxidation process. The formation of the Y2 Si2O7 phase further reduced the oxidation rate, as indicated by Fig. 4.  \\x0c', 'J. He et al. / Journal of the European Ceramic Society 36 (2016) 3769-3774   3773  Doping with aluminium was found to reduce the silica activity. However, as the temperature increased, the effect of the aluminium is reduced. As reported by Dhima et al. [22], at 1800  C, the activity of aluminium-doped silica is close to that of pure silica. This is why ZSA only exhibits a superior oxidation resistance than ZS at 1700  C. For the SiC with a low yttrium concentration (ZSY-2), the amount of yttrium in the ceramic composite is not high enough to fully stabilize the silica layer, thus the ZSY-2 exhibits an oxidation rate similar to the rate obtained for ZSA  in the  later stages of the oxidation process. The formation of Y2 Si2O7 in ZSY-7 indicates that enough yttrium  is dissolved  in  the silica, resulting  in  the  lowest in this study (10 times  oxidation rate observed  lower than the oxidation rate of the ZSA and ZSY-2 samples). The polished cross-sections of the samples oxidised at 1700  C for 60 min are compared in Fig. 6. Because of the severe oxidation of the ZS sample and the resulting oxide spallation, the cross-section of the ZS samples oxidized at 1700  C for 10 and 40 min are shown instead in Fig. 6(a-1) and (a-2). In Fig. 6(a-1), there is no silica layer on  the  top. Only a Zr-Si-O  layer and a broad SiC-depleted  layer were formed. Further oxidation  led to a deficiency of silica  in the Zr-Si-O  layer and  its eventual spallation, exposing  the depleted layer on the top (Fig. 6(a-2)).  It  is evident that the silica becomes volatile at 1700  C, which is in agreement with the results discussed above. The ZSA sample shows a triple-layered structure: a very thin SiO2 outer layer, a SiO2 layer with ZrO2 particulates dispersed in the Zr-Si-O layer, and a relatively thick SiC-depleted layer adjacent to the unaffected ZrB2 -SiC substrate (Fig. 6(b)). The oxide scales found on the ZSY-2 and ZSY-7 samples are different from the scales found on ZSA in the following two aspects (cp. Fig. 6(c) and (d)): First, the oxide scale contains only two layers: a Zr-Si-O outer layer and a SiCdepleted layer. The usual silica layer could not be detected, probably because it was too thin to be observed. Crystalline zirconia with a parallel column shape was found to be growing almost vertical to the oxidation surfaces, while the SiO2 fills the space between the zirconia grains. Second, the thickness of the SiC-depleted layer was thinner compared with the ZSA and ZS samples. Furthermore, the higher the yttrium concentration in SiC, the thinner the thickness of the SiC-depleted layer. However, no yttrium was detected by EDS in the samples probably because the concentration of the dopants in the samples is below the observation threshold. The oxidation of ZrB2 -SiC ceramics originates  from the active oxidation of SiC, which has been well documented by previous studies [1-3,7,11]. The active oxidation was found to be enhanced at elevated temperatures.  In order to  inhibit the active oxidation by further lowering the oxidation rate, a reduction of the activity of the silica layer is a promising approach. The addition of aluminium into SiC, which was reported to significantly increase the oxidation resistance at 1500  C,  is a good example  [12,22]. Although doping SiC with aluminium still resulted  in a  lower oxidation rate at 1700  C compared with undoped ZS, the effectivity of such a doping process was found to be very low. By replacing aluminium with yttrium, which has a larger atomic radius, the active oxidation of SiC could be further inhibited, as confirmed by the thinner SiC-depleted layer (Table 1).  4. Conclusions  In this study, ZS, ZSA, ZSY-2 and ZSY-7 samples were oxidised at 1700  C  for up to 1 h to  investigate the effect of a doping with Al or Y on the oxidation behaviour. The results showed that all the samples—except ZS—exhibited parabolic oxidation kinetics. For the ZS samples, the volatility of silica at 1700  C was found to be the reason for its rapid oxidation and linear oxidation kinetics. Doping SiC with Al/Y greatly improved the oxidation resistance of ZrB2 -SiC the silica activity at 1700  C. The ceramics due  to a reduction of   initial yttrium diffusion  towards  the surface of  the ZSY samples during the oxidation results  in a  fast oxidation during the  initial stage. However, the oxidation rate reaches a very  low  level once the yttrium concentration at the surface exceeds a certain value.  Acknowledgements  This work was financially supported by the Chinese Natural Science Foundation  (Grant Nos. 51172181 and 51521061), and  the “111” Project (Grant No. B08040).  References  [6]   [9]   [13]   [1] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Oxidation of zirconium diboride-silicon carbide in CO-CO2 at 1500   C at a low partial pressure of oxygen, J. Am. Ceram. Soc. 89 (10) (2006) 3240--3245. [2] M.M. Opeka, I.G. Talmy, J.A. 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He et al. / Journal of the European Ceramic Society 36 (2016) 3769-3774  [26]   I. Satoru, N. Toshiyuki, T. Hidehiko, Precipitation processing to synthesize fine polycarbosilane particles for precursors of silicon carbide powders, J. Ceram. Soc. Jpn. 114 (1330) (2006) 507-510. [27] Y. Wang, Y. Sohn, L. An, Y. Fang, L. Zhang, Oxygen transport through aluminum-containing amorphous SiO2 , J. Phase Equilib. Diffus. 27 (6) (2016).  [28] Y. Wang, Y. Fan, L. Zhang, W. Zhang, L. An, Oxidation behavior of SiAlCN ceramics at 1400   C, Scr. Mater. 55 (2006) 295-297. [29] K. Komeya, Y. Haruna, T. Meguro, Oxidation behaviour of Si3N4 -Y2O3 -Al2O3 system the sintered, J. Mater. Sci. 27 (1992) 5727-5734.  \\x0c']"
},{
  "_id": 162,
  "PDF": "Oxidation behaviours of ZrB2-SiC-MoSi2 composites at 1800 °C in air with different pressures.pdf",
  "Text": "['Corrosion Science 157 (2019) 87-97  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Oxidation behaviours of ZrB2-SiC-MoSi2 composites at 1800 °C in air with different pressures  T  Yang Yanga,b, Meishuan Lia, Lin Xuc,⁎⁎, Jingjun Xua,⁎, Yuhai Qiana, Jun Zuoa, Tongqi Lic  a Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China b School of Materials Science and Engineering, University of Science and Technology of China, Shenyang 110016, China c Science and Technology of Advanced Functional Composite Laboratory, Aerospace Research Institute of Materials & Processing Technology, Beijing 100076, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Oxidation behaviour ZrB2-SiC composite Low pressure air MoSi2  1.  Introduction  The oxidation behaviours of ZrB2-20SiC and ZrB2-20SiC-10MoSi2 (vol.%) at 1800 °C in air with different pressures of 0.5, 5 and 10 kPa (correspondingly the oxygen partial pressure was 0.1, 1 and 2 kPa, respectively) were investigated. In 0.5 kPa air, catastrophic oxidation occurred for both composites. However, the addition of MoSi2 could enhance the oxidation resistance of the ZrB2-20SiC composite in 5 and 10 kPa air, this was attributed to the reactions of MoSi2 with ZrB2 and highly volatile B2O3 to form stable MoB, causing the formation of a continuous, compact and silica-rich outer oxide layer with lower oxygen permeability.  Ultra-high temperature ceramics (UHTCs) possess the great potential to be applied in various extreme environments, for example, hypersonic aerospace vehicles, atmospheric re-entry and rocket propulsion system [1-3], due to their unique combination of properties, such as ultra-high melting point, excellent high temperature mechanical properties, and thermal as well as chemical stability [4]. For the application of UHTCs, the good oxidation resistance is an essential issue. As a member of UHTCs, ZrB2-SiC-based composites receive more and more attention due to their favourable oxidation resistance. A large amount of oxidation/ablation tests, conducted at 1000-2200 °C by using conventional furnaces in 1 atm static air [5-8] as well as plasma wind tunnel, oxyacetylene torch or arc-jet in dynamic atmosphere [9-11], demonstrated that a stable high viscosity silica-rich glass forms on the top surfaces of ZrB2-SiC composites, which provide the excellent oxidation resistance [12,13]. However, some recent investigations also showed that ZrB2-SiC composites exhibited significantly unusual oxidation behaviour at 1500 °C in 15 kPa O2/Ar(or N2) mixture gas with low oxygen partial pressure (pO2 ). For example, the protective silicarich outer layer disappeared and the oxide scale became porous and thick, consequently, a oxidation kinetics transition occurred from diffusion rate-controlled (parabolic law) one to reaction rate-controlled (linear law) one [14]. Such obvious discrepancy of oxidation performance of ZrB2-SiC composites mainly derived from the difference of environmental atmospheres, such as oxygen partial pressure [15], total  gas pressure and gas flow rate. As we know, the oxygen partial pressure is one of the most predominant environment factors affecting the oxidation behaviours of materials. For ZrB2-SiC composites containing C and B species, a great number of gaseous products, such as B2O3(g), CO (g) and CO2(g), formed during ultra-high temperature oxidation. Total gas pressure and gas flow rate could affect both the outward escape of gaseous products and the inward transport of oxygen, finally resulting in the change of oxidation behaviour. Considering the service condition of vehicles in extreme environments, low pressure air is usually involved. Therefore, it is necessary to explore the effects of air pressure (i.e oxygen partial pressure) on the oxidation resistance of ZrB2-SiC composites. Unfortunately, such relevant researches are very limited. Gao et al. investigated the oxidation behaviour of ZrB2-SiC in 20 kPa O2/N2 gas with pO2 = 200 Pa, found the formation of zircon was promoted below 1600 °C, which was thought to be favourable to the oxidation resistance [16]. Seong et al. and Han et al. reported that increasing SiC content was deleterious for the oxidation resistance of ZrB2-SiC composites under the condition of low oxygen partial pressure [17,18]. More recently, Jin et al. found that the oxidation kinetic of ZrB2-SiC-C composite at 1800 °C changed from parabolic to linear with decreasing pO2 [19]. However, these experiments were conducted either at temperatures below 1800 °C or in fixed pressure gases. The effect of total air pressure on oxidation behaviour of ZrB2-SiC composites as well as their oxidation mechanism in low pressure air at ultra-high temperatures were rarely discussed. In addition, it has been well known that, low oxygen  partial  ⁎ Corresponding author at: No. 72, Wenhua Road, Shenghe District, Shenyang 110016, China. ⁎⁎ Corresponding author at: No. 1, South Dahongmeng Road, Fengtai District, Beijing 100076, China. E-mail addresses: xulinhit@126.com (L. Xu), jjxu@imr.ac.cn (J. Xu).  https://doi.org/10.1016/j.corsci.2019.05.027 Received 7 March 2019; Received in revised form 27 May 2019; Accepted 29 May 2019  Available online 30 May 2019 0010-938X/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'Y. Yang, et al.  pressure will trigger the active oxidation of SiC, and it becomes even more severe with the increase of temperature [20]. It was the active oxidation of SiC that resulted in the absence of silica-rich outer layer and the change of oxidation kinetics for ZrB2-SiC composites in hightemperature oxygen-lean environment [21]. On the other hand, the active oxidation of SiC could also be triggered beneath the protective silica-rich layer, where the low oxygen partial pressure was induced by low diffusion rate of oxygen through the outer amorphous layer [22]. Consequently, a SiC-depleted layer was frequently observed near the unreacted substrate [6,7,9,16,19]. Which is consistent with a common oxidation phenomenon in alloys, i.e. internal oxidation, in which oxygen diffuses into the alloy substrate and causes the oxidation of one or more alloying element(s) internally [23]. Similarly, the formation of SiC-depleted layer underneath the oxide scale could be regarded as the internal oxidation of ZrB2-SiC composites. Commonly, this phenomenon will result in a porous layer between unaffected substrate and oxide scale, which deteriorates both of the oxidation resistance and mechanical properties. Therefore, to further and fully understand the oxidation resistance and evolution process of ZrB2-SiC composites, it is urgent to explore the influence of air pressure on their oxidation behaviours at ultra-high temperatures. In our previous work, the composites of ZrB2-20SiC and ZrB2-20SiC10MoSi2 (vol.%) were prepared by hot pressing, and it was found the that room-temperature mechanical properties of ZrB2-20SiC composite and oxidation resistance at 1800 °C in 1 atm air were improved significantly due to MoSi2 addition [24]. In this paper, the oxidation behaviours of ZrB2-20SiC and ZrB2-20SiC-10MoSi2 (vol.%) composites were further investigated at 1800 °C in air gases with different pressures, i.e. 0.5 kPa, 5 kPa and 10 kPa. The oxidation-induced mass change of two composites was determined, and the phase compositions as well as morphologies of oxidation zone were characterized. Finally, the oxidation process of ZrB2-20SiC composite in low pressure air and the effect of MoSi2 addition on its oxidation behaviour were discussed.  2. Experimental procedure  2.1. Preparation of  two ZrB2-based composites  Two ZrB2-based UHTCs, ZrB2-20SiC (labelled as ZS, vol.%,) and ZrB2-20SiC-10MoSi2 (labelled as ZSM) were synthesized by using hot pressing method. The commercially available raw powders of ZrB2 (99.5% purity, 1-2 μm size, Beijing HWRK Chem Co. LTD., Beijing, China), SiC (99% purity, 2 μm size, Forsmam (Beijing) Technology Co. (99.9% purity, 1-3 μm size, Beijing LTD., Beijing, China) and MoSi2 HWRK Chem Co. LTD., Beijing, China) were used for sample preparation. Two batches of powder mixture, with the values of 80:20:0 and 70:20:10 for the volume ratio of ZrB2:SiC:MoSi2, were planetary ballmilled for 12 h and then synthesized via hot-press sintering in a BNcoated graphite die under flowing argon atmosphere at 1900 °C. The details of the synthesis process were described elsewhere [24]. The densities of two as-sintered composites were measured by Archimedes method in deionized water.  2.2. Oxidation testing  Square sheets with the dimensions of 8 mm × 8 mm × 2 mm were diced from the hot pressed billets. After being chamfered, ground and finally polished with 0.5 μm diamond polishing slurry, all samples were ultrasonically cleaned in acetone and dried by flowing hot air. Oxidation test was conducted in an induction-heating ultra-high temperature oxidation testing apparatus which was built in our laboratory [25]. The sheet sample was placed on a graphite susceptor positioned in the centre of a water-cooled hollow copper induction coil. Then the ambient air pressure of the testing chamber was set as 0.5 kPa, 5 kPa and 10 kPa (correspondingly, the oxygen partial pressure was 0.1, 1 and 2 kPa), via a pressure control unit (ZCMD-I-10-LED, Chengduzhonghua,  Corrosion Science 157 (2019) 87-97  China, 100-10,000 Pa). A stable air flow was introduced into the chamber to maintain the constant chamber pressure. After the chamber pressure was stabilized to the desired value, the sample was heated, its surface temperature reached 1800 °C in about 2 min and was hold for 30 min. After that the sample was freely cooled to room temperature in about 20 min. Such rapid heating and cooling rates can simulate the high aerodynamic heating during hypersonic flight and minimize the effects of these two stages. Sample surface temperature was monitored by a two colour ratio pyrometer (Marathon MR1SCSF, Raytek, U.S.A., its accuracy: ± (0.5% Tmeas +2 °C), Tmeas in °C; measurement range: 1000-3000 °C). During cooling, the air pressure in the chamber always remained the same as that during isothermal testing. The exposed surface area (A) of the tested sample was measured prior to oxidation. The sample was weighed before (w0) and after (w1) oxidation test using an electronic balance (BP211, Sartorius, Germany) with an accuracy of 10−5 g. The weight change per unit surface area (Δwc) of the sample after oxidation was determined by:  Δ c w  =  w  1  w  0  −  A  (1)  2.3. Component and structure characterization  Kα  X-ray diffractometer (Cu radiation, D8 Advance, Bruker, Germany) was used to identify crystalline oxidation products. The surface and cross section morphologies of the thermally grown oxide scale were observed using a SUPRA 35 scanning electron microscope (SEM, LEO, Oberkochen, Germany) equipped with an energy-dispersive spectroscopy (EDS Inca X-Max, Oxford Instrument, Oxford, U.K.), and the light elements were characterized by EPMA (EPMA-1610, SHIMADZU, Japan). For cross-sectional analysis, the as-oxidized samples were firstly mechanical cut using a diamond slicing blade (Isomet 5000, Buehler, Lake Bluff, IL), then ground and polished after being mounted in resin. Before SEM observation, a thin layer of Au was sputtered on the surface of as-oxidized sample and the cross section of mounted sample to enhance their surface electric conductivity.  3. Result  3.1. Characterization of oxidation products  XRD patterns of the oxide scales formed on ZS and ZSM samples after oxidation at 1800 °C for 30 min in different pressure air are shown in Fig. 1, and the identified crystalline phases for the oxide scales are listed in Table 1. For both ZS and ZSM composites, ZrO2 with main monoclinic and minor tetragonal form were found under all conditions, while ZrSiO4 appeared only in 0.5 and 5 kPa air. The latter phenomenon indicates that low oxygen partial pressure favours the formation of zircon, which is consistent with the result obtained by Gao et al. [16]. Especially, for the ZSM composites, MoB was identified in all conditions, and crystal SiO2 and matrix ZrB2 were also detected only in 10 kPa air. Moreover, in all cases, no boron-containing crystalline oxide phase was detected by XRD in the oxide scales, indicating that the boron-containing oxide might escape into the surrounding environment as volatile gas or exist in the form of amorphous phase.  3.2. Oxidation-induced weight change  The weight changes of the two composites after oxidation at 1800 °C for 30 min in 0.5, 5, 10 kPa air are listed in Table 2. It can be seen that, for the ZS composite, its weight loss in 0.5 kPa air was −8.51 mg cm−2, small weight gain of 0.74 mg∙cm-2 in 5 kPa air was obtained, and less weight loss of −1.21 mg cm-2 was found in 10 kPa air. While for the ZSM composite, its weight change monotonously increased from obvious weight loss of −12.79 mg cm−2 in 0.5 kPa air to slight weight gain of 0.67 mg cm-2 in 10 kPa air. These results indicate that in low  88  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  Fig. 1. XRD patterns of  the oxide scales formed on (a) ZS and (b) ZSM composites after oxidation at 1800 °C for 30 min.  pressure air, a great amount of volatile products formed during ultrahigh temperature oxidation, because it was determined previously that the weight gain of these two ZS and ZSM composites was 3.06 mg cm−2 and 3.59 mg cm−2 after oxidation at 1800 °C for 30 min in 1 atm air, respectively [24].  3.3. Morphology of oxide scale  3.3.1. Surface morphology The macrographs of two composite samples oxidized at 1800 °C for 30 min in air under different pressures are shown in Fig. 2. It can be seen that with the increase of air pressure, the colour of the as-oxidized ZS sample changed from white to greyish white, and some spotted glass appeared on the sample surface after exposure in 10 kPa air. While the colour of the as-oxidized ZSM sample changed from greyish to partially dark grey and then to completely dark grey, suggesting that the outer glassy layer became continuous gradually. Meanwhile, bubbles appeared on the surface of ZSM sample after oxidation in 10 kPa air. Moreover, glass bubbles were also found on the side surfaces of both ZS and ZSM samples oxidized in 5 kPa and 10 kPa air, indicating that a glassy layer fully covered these side surfaces. Fig. 3 displays surface morphologies of the ZS composite after oxidation at 1800 °C for 30 min in air with different pressure. In the case of 0.5 kPa air, numerous nodules with the size of 12-20 μm and pores with the size of 5-8 μm appeared on the flat surface (see the magnification in Fig. 3(a)), and absence of the amorphous species. A similar surface morphology was also observed on ZrB2-SiC-based composites after oxidation at low oxygen partial pressures [20,21,26]. In the case of 5 kPa air, the formed oxide scale was grey and dispersedly embedded white fine particles (˜1-3 μm). According to XRD and EDS analysis, these white fine particles were identified as ZrO2, while the grey oxide was mainly SiO2 glass. And in some local areas, only glassy SiO2 existed on the top surface. When the air pressure was 10 kPa, the surface morphology was similar to that in 5 kPa air, but the grey amorphous SiO2 was interconnected, accordingly the emerged ZrO2 particles decreased, as shown in Fig. 3(c). For the ZSM composite, after oxidation in 0.5 kPa air, its outer surface was also covered with a similar oxide scale, in which a great  ˜15 μm and number of nodules and pores, with the smaller sizes of ˜5 μm respectively, existed. And no amorphous species could be distinguished as well (see Fig. 4(a)). When the air pressure was raised to 5 kPa, the surface morphology of oxide scale was similar to that of ZS, i.e. fine ZrO2 particles was embedded in grey SiO2 glass. Meanwhile, more amorphous SiO2 existed and interconnected, as shown in Fig. 4(b). Grey SiO2 glass layer covered completely the ZSM sample surface after oxidation in 10 kPa air (see Fig. 4(c)). It is worth noting that some bubbles appeared on the outermost layer of ZSM sample after oxidation in 5 and 10 kPa air, as well as of ZS sample in 10 kPa air. Such bubbles were also observed on the surfaces of ZrB2-SiC-based composites after oxidation at temperature above 1500 ℃ in air or wind tunnel test [7-9,15,27].  3.3.2. Cross section morphology The cross section views and corresponding EDS mappings of O, Si, Zr for ZS composite after oxidation at 1800 °C for 30 min in air with different pressures are revealed in Fig. 5. It can be seen from Fig. 5(a) that after oxidation in 0.5 kPa air, a single porous oxide scale with 130 μm thick formed, which was composed of mainly zirconia and minor Si-containing oxide (ZrSiO4). Besides, a SiC-depleted layer with thickness of about 50 μm existed between the oxide scale and unreacted substrate. In 5 kPa air, the oxide scale could be divided into thin SiO2rich outer layer and relatively thick highly porous ZrO2-rich layer. Expectedly, a SiC-depleted layer with thickness of about 24 μm also existed between the oxide scale and unreacted substrate (see Fig. 5(b)). After oxidation in 10 kPa air, the oxide morphology was quite similar to that in 5 kPa air, but the SiO2-rich outer layer and SiC-depleted layer became thicker, as shown in Fig. 5(c). Undoubtedly, the internal oxidation occurred for the ZS composite during oxidation at 1800 °C in air with reduced pressure range of 0.5-10 kPa, as a SiC-depleted layer existed after oxidation under all test conditions. Fig. 6 shows the cross-sectional micrographs and corresponding EDS mappings of O, Si, Zr for ZSM composite after oxidation at 1800 °C for 30 min in air with different pressures. After oxidation in 0.5 kPa air, the reaction zone of ZSM composite was similar to that of ZS composite. The ZrO2-rich single-layer oxide scale was porous and its thickness was about 120 μm. A SiC-depleted layer with the thickness of 60 μm existed  Table 1  Identified oxidation products of ZS and ZSM composites via XRD and EDS.  Air pressure (kPa)  ZS  ZSM  0.5 5 10  m-ZrO2, t-ZrO2 l, ZrSiO4 m-ZrO2, t-ZrO2 l, ZrSiO4 m-ZrO2, t-ZrO2 l, SiO2  a  l  l, SiO2  a  Note:  l   little; a amorphous, undetected by XRD.  89  m-ZrO2, t-ZrO2 l, ZrSiO4 l, MoBl m-ZrO2, t-ZrO2 l, ZrSiO4 l, MoB, SiO2 m-ZrO2, t-ZrO2 l, SiO2, MoB, SiO2  a  a  \\x0c', 'Y. Yang, et al.  Table 2  Corrosion Science 157 (2019) 87-97  The weight change and oxide thickness of two composites after oxidation at 1800 °C for 30 min in air atmospheres with different pressures.  Sample ID  Air pressure (kPa) Weight change (mg cm−2) SiO2-rich layer (μm) ZrO2-rich layer (μm) SiC-depleted layer (μm)  ZS  0.5 −8.51 -  130.8 ± 3.9 50.4 ± 5.1  * Thickness of (SiO2+ZrO2) sublayer.  5 0.74 3.2 ± 2.3* 127.2 ± 2.5 24.5 ± 4.7  10 −1.21 4.8 ± 1.3 101.5 ± 4.5 34.5 ± 5.4  ZSM  0.5 −12.79 -  120.7 ± 3.7 59.9 ± 7.1  5 0.05 14.2 ± 0.8 57.4 ± 2.6 22.1 ± 2.2  10 0.67 4.4 ± 0.5 28.4 ± 2.5 -  between the oxide scale and unreacted substrate, as revealed in Fig. 6(a). After oxidation in 5 kPa air, the bi-layered oxide scale formed, consisted of a SiO2 outer layer with 14 μm thick and a ZrO2-rich inner layer (see Fig. 6(b)). A SiC-depleted layer was also observed in the substrate just beneath the oxide scale. Both the thicknesses of the oxide scale and SiC-depleted layer were thinner compared to that formed in 0.5 kPa air. After oxidation in 10 kPa air, the oxide scale was dense and thinner (˜36 μm), which also possessed bi-layered structure with a SiO2rich outer layer and a ZrO2-SiO2 mixture inner layer. No SiC-depleted layer was found. Moreover, the ZrO2-SiO2 layer had a sandwich-like structure with a semi-continuous silica layer (˜2 μm) inside. The similar structure of oxide scale formed on ZrB2-MoSi2 was also observed by Liu et al. [28]. In 5 and 10 kPa air, well-distributed bright particles with the size of ˜2 μm were found in ZrO2 layer. According to the results of EPMA, they were identified as MoB, which was also reported in the oxide scale of ZrB2-based UHTCs containing MoSi2 [6,27,28]. The thicknesses of the oxide scale and SiC-depleted layer for all samples are summarized in Table 2. It can be seen that the thickness of the oxide scales of ZS and ZSM composites shared a same trend with weight change as the increase of air pressure. For ZS composite, the oxide scale formed after oxidation in 5 kPa air was the thickest, while the SiC-depleted layer was the thinnest, compared with that formed in 0.5 and 10 kPa air. The thicknesses of the oxide scale and SiC-depleted layer of ZSM composite decreased monotonously with increasing air pressure. So, it was concluded that the total thickness of their reaction zones was comparable under the condition of 0.5 kPa air, but it was much thinner for ZSM composite in the cases of 5 and 10 kPa air.  4. Discussion  4.1. Characterization of oxidation resistance  In general, the oxidation resistance of metals can be simply evaluated via their apparent weight gain, as the formed oxides are condensed (mostly solid). However, during oxidation of UHTCs, gaseous and/or volatile products, such as CO for carbides, N2 for nitrides, B2O3 for borides, are generated. As a result, their oxidation resistance can`t be determined only by the apparent weight gain or weight loss due to the uncertainties in the quantity of escaped species. Usually, the thickness of oxide scale is taken as a criterion for the evaluation of oxidation resistance of these composites [5,16,19,26,28]. However, this method is questionable for the same reason, no considering the physical non-denseness of oxide scales due to the presence of pores, cracks, bubbles etc. So, in the present work, the oxidation resistance of the two composites was characterized via the consumption of substrate materials induced by the oxidation. It should be noted that both of the external oxidation and internal oxidation were involved during the oxidation of ZS and ZSM composites. For external oxidation, ZrO2, formed by the oxidation of ZrB2, was the only non-volatile species among all detected oxidation products. Therefore, the consumption of ZrB2 could be calculated according to the amount of ZrO2. Similarly, based on the nominal compositions of ZS and ZSM, the substrate consumption caused by external oxidation could be estimated for each sample from the consumption of ZrB2. For internal oxidation, only SiC in the substrate was oxidized preferentially to form SiO(g) and CO(g) gases, accompanied with the presence of a SiCdepleted layer. During internal oxidation, the ZrB2 matrix maintained its original structure, i.e. only voids were left in this layer after active oxidation of SiC. Consequently, the total volume (thickness) was’ t  Fig. 2. Macrographs of ZS and ZSM composites after oxidation at 1800 °C for 30 min in different pressure air.  90  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  Fig. 3. Surface morphologies of ZS composite after oxidation at 1800 °C for 30 min in 0.5 (a), 5 (b) and 10 (c) kPa air.  Fig. 4. Surface morphologies of ZSM composite after oxidation at 1800 °C for 30 min in 0.5 (a), 5 (b) and 10 (c) kPa air.  changed. Therefore, the substrate consumed by external oxidation was proportional to the thickness of the formed ZrO2-rich layer, as the change of the total surface area of tested samples was negligible during oxidation. And internal oxidation caused the appearance of the SiC-depleted layer. In this way, it was reasonable that the oxidation resistance of the composites can be characterized via ZrO2-rich layer thicknesses. It can be seen from Table 2, with decreasing of air pressure, the ZrO2rich layer increased for both of the composites. And the ZSM exhibited a much lower consumption in 5 and 10 kPa air than that of ZS, indicating that the consumption of ZrB2-SiC could be effectively retarded by the addition of MoSi2 in 5 and 10 kPa air. While for both ZS and ZSM, the thicknesses of the non-protective oxide scale and the internal oxidation zone were massive under the oxidation condition in 0.5 kPa air. The following sections will discuss the exact mechanism behind these phenomena thermodynamically.  4.2. Oxidation process of ZrB2-20SiC  4.2.1. External oxidation According to the results of XRD and SEM, the major reactions expected to take place in ZS composite during the oxidation are listed below:  2/5ZrB2(s) + O2(g) → 2/5ZrO2(s) + 2/5B2O3(l)  B2O3(l) → B2O3(g)  2/3SiC(s) + O2(g) → 2/3SiO2(l) + 2/3CO(g)  SiO2(l) → SiO2(g)  (R1)  (R2)  (R3)  (R4)  As described by the reactions of R1 and R3, during the oxidation at 1800 °C in air, the main oxidation products of ZrB2 are ZrO2 and B2O3, and those of SiC are SiO2 and CO. Sequentially a protective oxide scale with a structure of “solid pillars with liquid roof” was formed, in which the borosilicate glass acted as a barrier for inward diffusion of oxygen, meanwhile ZrO2 grains retained the glass phase and provided mechanical stability to the oxide layer [13]. However, the evaporation of  91  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  Fig. 5. The cross sections and corresponding EDS mappings of O, Si, Zr for ZS composite after oxidation at 1800 °C for 30 min in 0.5 (a), 5 (b) and 10 (c) kPa air.  volatile species was inevitable at 1800 °C [25], especially under low air pressure. Therefore, B2O3(l) and SiO2(l) in the glassy roof would volatilize via reactions of R2 and R4 under the testing condition, respectively. To better estimate the vapour pressure of these volatile species under the testing conditions, the volatility diagram of ZrB2-SiC at 1800 °C was compiled based on the data from NIST-JANAF tables [29], as shown in Fig. 7. During the calculation, several simplified assumptions were referred to the model established by Fahrenholtz [22,30]. For simplification, only the major vapour phases (volatile species with the highest equilibrium vapour pressure) were considered, because minor volatile species had negligible practical consequences to the system. While, as the phases of interest, the equilibrium vapour pressures of SiO (g) and SiO2 (g) were also shown in the diagram. For convenience, various oxygen partial pressures applied during tests were marked with dash lines in the diagram, i.e. pO2 = 0.1, 1 and 2 kPa, correspondingly the total air pressure of 0.5, 5 and 10 kPa, respectively. It can be seen from Fig. 7, B2O3 (g) was the predominant volatile gas  with a constant vapour pressure of ˜7.14 kPa under all test conditions. The vapour pressure of B2O3 (g) approached to 10 kPa, much higher than the air pressures of 0.5 kPa and 5 kPa. The vapour pressure of SiO (g) decreased as oxygen partial pressure increased, while pSiO2 main˜0.08 Pa. Normally, weight tained a constant of loss induced via evaporation at high temperature could be measured only if the vapour pressure of the volatile species was higher than 0.001 Pa [31,32]. Apparently, the vapour pressures of all these species were much higher than 0.001 Pa, indicating that vigorous volatilization occurred during all tests (especially, boron oxide would volatilize immediately once it formed). And the escape of SiO(g) became even more serious under lower oxygen partial pressure, resulting in the reduced protective SiO2, as shown by the surface and cross-sectional micrographs of ZS composite in Fig. 5. Moreover, at 1800 °C, the vigorous gaseous volatilization of B2O3 was inevitable for ZS composite. Therefore, the single porous ZrO2-rich layer formed, and no integrated SiO2-rich outer layer could exist even in 5 and 10 kPa air where SiC was oxidized passively, as the formation and rupture of these gaseous bubbles would destroy  92  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  Fig. 6. The cross sections and corresponding EDS mappings of O, Si, Zr for ZSM composite after oxidation at 1800 °C for 30 min in 0.5 (a), 5 (b) and 10 (c) kPa air.  the compactness and completeness of the protective “liquid roof”.  4.2.2. Internal oxidation—active oxidation of SiC As mentioned above, internal oxidation, i.e. active oxidation of SiC (reaction R5), was triggered in the substrate adjacent to the oxide scale, where oxygen partial pressure became low enough. The formed SiO (g) escaped outwardly from the internal oxidation zone. When it escaped to the position, where the oxygen partial pressure was higher enough, it could be further condensed into SiO2 (l), as denoted by R6 [25]. The top surface of oxide scale has the highest oxygen pressure (equals to ambient). So, an oxygen partial pressure gradient was established, and the internal oxidation process of ZrB2-SiC was proposed schematically, as shown in Fig. 8. At the upper side and bottom of the internal oxidation zone, the reactions R1 and R5 occurred. Fahrenholtz calculated the oxygen partial pressure of SiC-depleted layer boundaries at 1500 °C [22]. Based on his model, the determined equilibrium oxygen pressures of R1 and R5 were 3.6 × 10−8 and 4.2 × 10−10 Pa at 1800 °C, respectively. The reactions and corresponding equilibrium oxygen  93  Fig. 7. The volatility diagram of ZrB2-SiC composite at 1800 °C.  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  Fig. 8. Schematic illustration of the formation of complex reaction zones of ZS (a) and ZSM (b) composites during oxidation and various interface reactions under the oxidation conditions at 1800 ℃ in 0.5 and 5, 10 kPa air.  pressures are also provided in Fig. 8.  SiC(s) + O2(g) → SiO(g) + CO(g)  SiO(g) + 1/2O2(g) → SiO2(l)  only take place underneath the oxide scale, i.e. in the external oxidation zone. And SiO(g) could be further oxidized into SiO2(l) as a barrier to inhibit oxygen inward penetration.  (R5)  (R6)  On the other hand, the transformation from passive to active oxidation for SiC occurs, once the oxygen partial pressure is lower than a critical value. Generally, the critical pO2 increases with the increase of ˜230 Pa at 1800 ℃ temperature. And the critical value is as high as [33,34]. In this case, the initial oxidation of SiC for the two composites at 1800 ℃ in 0.5 kPa air (pO2 = 100 Pa) was active one, and two kinds of gaseous products formed (see R5), consequently no SiO2 existed on the oxide scale. Moreover, because of the outward transportation of massive gaseous products through the oxide scales, numerous pores appeared in the oxygen affected zone formed on the two composites in 0.5 kPa air, as shown in Figs. 3(a) and 4 (a). When the air pressure was increased to 5 and 10 kPa (pO2 = 1 and 2 kPa, respectively), passive oxidation of SiC occurred once exposed to such oxidizing environment. While, the active oxidation of SiC could  4.2.3. Oxidation induced weight change The weight change of the tested composite was caused by both external and internal oxidations. The oxidation reactions in these two zones were quite different under different oxidation conditions (i.e. air pressure), consequently, the weight change varied. The molar weight change in external and internal oxidation zones are theoretically calculated for ZS and ZSM materials, as shown in Table 3. The internal oxidation causes the same weight loss (over 11%) for all conditions, as no condensed products were produced. However, a transformation from weight loss to weight gain occurred when the air pressure increased from 0.5 kPa to 5-10 kPa for the case of external oxidation, as the formed SiO2 via SiC was partially remained in the oxide scale. Therefore, obvious weight loss occurred for ZS composite because both the external and internal zones were very thick (see  94  \\x0c', 'Y. Yang, et al.  Table 3  Corrosion Science 157 (2019) 87-97  The calculated weight change induced by various reactions in external and internal oxidation zones of ZS and ZSM composites.  composites  ZS  ZSM  Air pressure  Oxidation zone  Major reaction  Weight change (%∙mol−1)  Major reaction  Weight change (%∙mol−1)  0.5 kPa  5, 10 kPa  External Internal External Internal  (1), (2), (5) (5) (1), (2), (3), (4) (5)  −3.5 −11.7 +0.8˜+13.95 −11.7  (1), (2), (4), (5), (9) (5) (1), (2), (3), (4), (7), (8) (5)  −15.9˜-7.9 −11.6 −1.3˜+18.5 −11.6  Fig. 5(a)) in 0.5 kPa air. And, the weight loss reduced in the case of higher air pressure, even the small weight gain occurred.  4.3. Oxidation of ZrB2-20SiC-10MoSi2  According to XRD, SEM and EPMA analysis results, MoB was identified in the oxide scales of ZSM composite under all test conditions (see Table 1 and Fig. 6). However, no other oxidation products of MoSi2, such as Mo5Si3 [35] and MoO3 [27], were detected. The absence of MoO3 could be due to its volatilization as gaseous phase. It is proposed that MoB was formed via the reactions between MoSi2 and ZrB2 or B2O3 (see R7 and R8). To thermodynamically determine the possibility of R1, R3, R7 and R8 at 1800 °C, the variations of standard Gibbs free energy ( m ) and tendency of equilibrium oxygen pressure versus temperature for these reactions were plotted in Fig. 9(a) and (b), respectively. It should be noted that the reaction R9 was also involved in the diagrams. And the calculation of m was performed based on unit molar oxygen consumption.  Δ Gr  Δ Gr  °  °  1/5ZrB2(s) + 2/5MoSi2(s) + O2(g) → 1/5ZrO2(s) + 2/5MoB(s) + 4/ 5SiO2 (l) (R7)  4/5MoSi2(s) + 2/5B2O3(l) + O2(g) → 4/5MoB(s) + 8/5SiO2(l)  (R8)  2/7MoSi2(s) + O2(g) → 2/7MoO3(g) + 4/7SiO2(l)  (R9)  °  °  Δ Gr  °  Δ Gr  Δ Gr  Fig. 9(a) shows that the values of m for the reactions of R7 and R8 are lower than that of R1 in the temperature range involved. Especially the m of R8 is even lower than that for R3 at temperature above 1740 °C, suggesting that the thermodynamic motivation of R8 is the highest among all these oxidation reactions at 1800 °C. On the other hand, m for the formation of MoO3 (R9) is much higher than that of the others, indicating that under the oxidation conditions, MoSi2 would react with B-containing compounds, such as ZrB2 and B2O3, preferentially, to form MoB (its melting point is above 2180 °C [36]). This is consistent to the results of XRD and EPMA analysis. Moreover, the equilibrium oxygen pressures [pO2 (R7) and pO2 (R8)] are also much lower than pO2 (R1) and pO2 (R9) (see Fig. 9(b)), suggesting that these  reactions (R7 and R8) could occur at much lower oxygen partial pressure even beneath the silica-rich layer. Therefore, in the ZSM composite, ZrB2 matrix together with MoSi2 was preferentially oxidized to produce MoB rather than B2O3 at 1800 °C. Moreover, B2O3, generated by the oxidation of ZrB2 matrix, together with MoSi2 could be reacted to form stable MoB (see R8). In this case, the formation of highly volatile B2O3 was inhibited, helpful to maintain the integrity of the glass layer. Moreover, R7 and R8 might occur at the substrate/oxide interface, where the pO2 was much lower than that on the top surface, and SiO2 was produced. As these reactions proceeded, the accumulated SiO2 at the interface resulted in the formation of an inner silica-rich layer, as shown in Fig. 6(c). The continuous SiO2 layer with a very low vapour pressure, is crucial for the oxidation resistance of this composite. All of these factors resulted in the improvement of oxidation resistance of ZSM sample at low oxygen partial pressure, finally a thinner oxide scale existed compared with ZS sample. Moreover, the addition of MoSi2 also affected the internal oxidation of ZrB2-SiC composite, because the internal oxidation zone of ZSM composite was much thinner, as shown in Fig. 6. As mentioned above, the upper and lower boundaries of pO2 in this region formed in ZrB2-SiC composite were controlled by the equilibrium oxygen pressures of R1 and R4, respectively. Correspondingly, for ZSM composite these boundaries was determined via the equilibrium oxygen pressure of R7 the upper pO2 and R4, respectively, as shown in Fig. 8. Therefore, boundary of the SiC-depleted layer in ZSM composite decreased from ˜3.6 × 10−8 Pa to ˜1.0 × 10−8 Pa at 1800 °C. Thus, the oxygen partial pressure range in the SiC-depleted zone was narrowed after the addition of MoSi2 in ZS material, indicating that the active oxidation of SiC at the substrate/oxide scale interface was also inhibited. Consequently, the thickness of the SiC-depleted layer was reduced for ZSM composite after oxidation in 5 and 10 kPa air. For the case of 0.5 kPa air, catastrophic oxidation also occurred for ZSM sample. Because the oxygen partial pressure was lower than 230 Pa, and initial active oxidation of SiC resulted in the absence of SiO2 outer-layer. Although, the vigorous volatilization of B2O3 was partly inhibited as the formation of MoB, other gaseous species, such as  Fig. 9. Thermodynamically analysis of the oxidation reactions R1, R3, R7, R8, R9 of ZSM composite in the temperature range of 1700-1900 °C, (a) Gibbs free energy ( m ), (b) equilibrium oxygen pressure.  Δ Gr  °  95  \\x0c', 'Y. Yang, et al.  Corrosion Science 157 (2019) 87-97  CO(g) and SiO(g), were still massive in 0.5 kPa air. Therefore, a porous oxide scale with numerous through-thickness pores formed on the surface, which is non-protective. It can be seen from Table 3, both external and internal oxidation for ZSM composite caused massive weight loss even higher than that of ZS composite in 0.5 kPa air. As a result, the weight loss of ZSM composite was much higher in this condition. According to Table 2, the ZrO2-rich layer thickness of ZSM composite is comparable to that of ZS composite, indicating the substrate consumption of these two materials are similar in 0.5 kPa air, as well as the oxidation resistance. However, in 5 and 10 kPa air, the theoretical molar weight change for external oxidation is -1.3˜+18.5% and that for internal oxidation stays constant. Thus, slightly weight gain occurred in 5 and 10 kPa for ZSM composite. What’s more, the oxidation zone thickness (especially the internal oxidation zone) of ZSM composite decreased remarkably with the increase of air pressure, indicating that the weight loss would highly decrease. The thinner ZrO2-rich layer corresponded a slight substrate consumption for ZSM composite, and more SiO2 existed in the outer layer. All of these distinctions indicated a better oxidation resistance for the ZSM material.  5. Conclusion  The oxidation behaviour of hot-pressed ZrB2-20SiC and ZrB2-20SiC10MoSi2 (vol.%) composites were investigated at 1800 °C in air with the pressure of 0.5, 5 and 10 kPa (the corresponding oxygen partial pressure was 0.1, 1 and 2 kPa, respectively). Both the composites exhibited catastrophic oxidation at pO2 = 100 Pa, and a thick non-protective porous ZrO2 layer formed after oxidation for 30 min, as the oxygen partial pressure was lower than the critical value (230 Pa) and intrinsic active oxidation of SiC under such condition. With increasing oxygen partial pressure, the oxide scale became thinner, but no continuous protective glass layer formed on ZrB2-20SiC composite, because of the vigorous evaporation of gaseous species, especially B2O3(g) and the existence of porous internal oxidation layer. On the contrary, the addition of MoSi2 resulted in the formation of a continuous silica-rich oxide layer, the oxidation affected zone was remarkably reduced, implying the improvement of oxidation resistance of ZrB2-20SiC-10MoSi2 composite. During the oxidation of ZSM composite, the formation of MoB via reaction between MoSi2 and ZrB2 inhibited the formation of volatile B2O3 and led to the presence of a continuous SiO2 outer layer in the oxide scale.  Acknowledgments  This work is financially supported by National Natural Science Foundation of china (Grant Nos. 51571203 and 51571205) as well as science and technology foundation of national defence key laboratory (Grant No. HTKJ2019KL703006).  References  [1] W.G. 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  "_id": 163,
  "PDF": "Oxidation mechanism and resistance of ZrB2–SiC composites.pdf",
  "Text": "['Corrosion Science 51 (2009) 2724-2732  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Oxidation mechanism and resistance of ZrB2-SiC composites  Ping Hu a,*, Wang Guolin b, Zhi Wang a  a Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, China b China Aerodynamic Research and Development Center, Mian Yang Si Chuan 621000, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 27 May 2009 Accepted 4 July 2009 Available online 9 July 2009  Keywords:  A. Ceramic C. Oxidation  1. Introduction  Transition metal borides and carbides such as ZrB2, ZrC, HfB2, HfC, and TaC have melting temperatures in excess of 3000 °C, making them candidates for use as structural materials at temperatures above 1800 °C [1]. Ultra-high temperature ceramics (UHTCs) have attracted great attention in recent years because of the drive to produce thermal protection systems (TPS) and other components for hypersonic aerospace vehicles [2-9]. Within the family of UHTCs, ZrB2 has the lowest theoretical density (6.09 g/ cm3) and a high thermal conductivity, which make it attractive for aerospace applications [8-10]. However, the oxidation of ZrB2 is very poor at temperatures above 1400 °C due to the volatilization of B2O3, which results in formation of a porous, non-protective ZrO2 layer [11,12]. Numerous investigations to improve the oxidation resistance of ZrB2 have been reported [5,11,13-18]. It was found that SiC addition provided the most oxidation resistance by promoting the formation of borosilicate glass. This borosilicate glass afforded more oxidation protection than boria since it is more viscous, has a higher melting temperature and a lower vapor pressure, and is more of a barrier to oxygen diffusion. Compositions with 5-50 vol.% SiC were investigated for ZrB2 over a wide range of test temperatures and pressures; 20 vol.% compositions were judged optimal for hypersonic vehicles in a series of efforts supported by the US Air Force [19,20]. The formation of a ‘‘SiC-depleted layer” during oxidation of ZrB2-SiC due to the active oxidation of SiC has been noted by several researchers [5,15-17]. A thermodynamic model has been developed to explain the formation of a SiC-depleted layer during ZrB2-SiC oxidation in air at 1500 °C [21]. However, the active to  * Corresponding author. Tel./fax: +86 45186402382. E-mail address: huping@hit.edu.cn (P. Hu).  0010-938X/$ see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2009.07.005  The oxidation mechanism of ZrB2-SiC composites was investigated based on a combination of theory and experiments. The oxidation reactions, microstructure evolution, scale stability and temperature limit were examined in our research and a good correspondence was obtained between theoretical predictions and experimental results. Microstructure evolution and stability are significantly dependent on both temperature and composition. SiO2 is thermochemically stable below 1800 °C and will lose its protective properties at temperatures above 2300 °C. The temperature limit for ZrB2-SiC composites is strongly dependent on the vapor pressure of the gaseous products and volume content of ZrB2. Ó 2009 Elsevier Ltd. All rights reserved.  passive transition in the oxidation of SiC and the effect of this layer on the oxidation resistance have not been recognized. The authors recently found that SiC-depleted layer formation in ZrB2-30 vol.% SiC significantly decreased the oxidation resistance of this composite at temperatures above 1800 °C [22,23]. In addition, cracking and spallation were observed in the SiC-depleted layer after oxidation in air at 1900 °C [23]. The oxidation resistance and mechanism are strongly related to the evolution of the oxide scale. The development of the layered structure on the surface of ZrB2-SiC air up to 1500 °C has been investigated in detail [24]. Very little attention has been paid to the structural evolution of the oxide scale at higher temperatures, especially at temperatures above 1800 °C. To understand oxidation behaviour of UHTCs, thermodynamic calculations have been conducted on the volatility diagrams [12,25]. Vapor pressure considerations provide significant insight into the relatively good oxidation resistance of ZrB2-based materials at 2000 °C and above. These materials form multi-oxide scales composed of a refractory crystalline oxide (skeleton) and a glass component, and this compositional approach is recommended for further development [25]. Unfortunately, the scale stability and temperature limit of these materials has not been quantitatively investigated. Up to now, the oxidation mechanism of UHTCs is still not well understood. The use of these materials in hypersonic and propulsion applications requires a more comprehensive understanding of the associated oxidation mechanism. In accordance with the previously reported results, a composite with a diboride matrix and a SiC as a second phase possesses one of the most promising compositions. Moreover, the properties of ZrB2-SiC composites can be modified or optimized for rigorous demands of ultra-high temperature environments. This paper describes the through a combination of  oxidation mechanism of theoretical analysis and  ZrB2-SiC selected  \\x0c', 'experiments. Better understanding of the oxidation mechanisms of ZrB2-SiC will provide a good scientific foundation and technical support for the revolutionary improvements in the design of aerospace vehicles and the capabilities of modern propulsion systems.  2. Theoretical considerations  The oxidation mechanism of ZrB2-SiC composites is significantly affected by the oxidation behaviours of SiC and ZrB2, respectively. It is well known that oxidation behaviour of SiC can be classified as active or passive oxidation. The important passive and active reactions for SiC are as follows:  SiCðcÞ þ 3 O2 ðgÞ ¼ COðgÞ þ SiO2 ðcÞ 2 SiCðcÞ þ O2 ðgÞ ¼ COðgÞ þ SiOðgÞ SiCðcÞ þ 1 2  ð1Þ  ð2Þ  O2 ðgÞ ¼ CðcÞ þ SiOðgÞ  ð3Þ  The most common route of SiO(g) generation is active oxidation. Although this can lead to rapid degradation, it only occurs at low oxygen potentials and high temperatures. The major issue with active oxidation of SiC is the PO2 transition point. For Reaction (2), the calculation of the oxygen pressure transition point is complex because two gaseous species are involved. Using Wagner’s model alone makes it hard to predict the active to passive transition in the oxidation of SiC. M. Balat et al. took both Eriksson’s model and Wagner’s model into account to determine the active to passive transition. The results are provided the following equation [26]:  Ptr O2 ¼  DSiO DO2  \\x12  \\x13  3=8  DCO DO2  \\x12  \\x13  1=8  K 3=4 2  K \\x001=2 1  ð4Þ  where, the ratios of DSiO =DO2 and DCO =DO2 can be assumed as 0.44 and 0.99, respectively. [26] The K1 and K2 are the equilibrium constant for Reactions (1) and (2), respectively. The calculation of active to passive transition for Reaction (3) can be obtained directly based on SiO equilibrium.  Ptr  O2 ¼ 0:5Peq  SiO  ð5Þ  The factor of 0.5 is from the stoichiometry of Eq. (3) and the PSiO for a stable SiO2 scale can be determined from the following equilibrium.  SiCðcÞ þ SiO2 ðlÞ ¼ CðcÞ þ 2SiOðgÞ  ð6Þ  The oxidation mode of SiC is strongly dependent on the temperature and oxygen partial pressure. Oxygen partial pressure versus temperature for the active to passive oxidation SiC under standard air is shown in Fig. 1. The calculations were performed using JANAF data [27]. The oxygen partial pressure for the active to passive oxidation is determined by Reactions (2) and (3), and the latter reaction is favored at lower oxygen partial pressures according to thermodynamic calculations. The transition temperatures for Reactions (2) and (3) are 1734 and 1931 °C in air, respectively, which decrease significantly with a reduction in oxygen partial pressure. At an oxygen partial pressure of 1 \\x02 104 Pa, the transition temperatures for Reactions (2) and (3) are 1665 and 1855 °C, respectively. The temperature for active to passive oxidation of SiC is only 1532 °C when the oxygen partial pressure decreased to 2 \\x02 103 Pa. It should be noted that the experimental values are always lower than the calculated values because the oxygen partial pressure at the reaction interface is lower than that in the ambient atmosphere and thus leads to a decrease of transition temperature. The total pressure at the reaction interface is a key factor in the mechanical stability of oxide scale. Calculations indicate that when T = 1809 °C the total pressure at the SiC/SiO2 interface reaches  1 atm, which is independent of oxygen partial pressure. Moreover, the peak equilibrium total vapor pressure inside the oxide scale could exceed 1 atm at temperatures above 1794 °C. Therefore, the formed oxide scale for SiC is mechanically unstable when the temperature is higher than 1794 °C resulting in active oxidation of SiC. ZrB2 oxidizes mainly according to Reaction (7). At high temperatures, B2O3(g) forms by the direct vaporization of B2O3(l) according to Eq. (8). Note that the oxidation of ZrB2 would produce the condensed phase (ZrO2) which is significantly different from the oxidation behaviour of SiC. Formation of the B-containing gases with high vapor pressure is an issue for the application of ZrB2 based composites. The B2O3 pressure approaches 1 atm for ZrB2 when the temperature is increased to 2066 °C, which is the upper use temperature of ZrB2 based composites.  ZrB2 ðcÞ þ 5 O2 ðgÞ ¼ ZrO2 ðcÞ þ B2O3 ðlÞ 2 B2O3 ðlÞ ¼ B2O3 ðgÞ  ð7Þ  ð8Þ  The volatility diagram for ZrB2-SiC at 2000 °C and PCO = 105 Pa is shown in Fig. 2. The predominant gaseous species for ZrB2-SiC at low and high oxygen partial pressures are SiO(g) and B2O3(g), respectively. A stable C(c) field is also present at low oxygen partial pressures, which is consistent with previous analyses of SiC active oxidation. The C(c) phase field will shrink with decreasing PCO. Moreover, C(c) is easily consumed by the following reaction:  CðcÞ þ 1  2  O2 ðgÞ ¼ COðgÞ  ð9Þ  1200  1500  1800  2100  2400  2700  0  1  2  3  4  5  6  7  Si C (c) + 1/2  O  2( g) =  C (c) +  Si O ( g)  Si C (c)+  O  2  (g)=  C  O (g)+  Si O (g)  C  O  +  Si O  SiC+SiO2  C+SiO  g o L  P  O  2  (  P  a  )  Temperature (  )  Fig. 1. Oxygen partial pressure oxidation SiC under standard air.  versus  temperature  for  the  active  to  passive  -12  -10  -8  -6  -4  -2  0  0  2  4  6  8  SiO(g)  B2O3(g)  ZrB2 (c)+SiO 2 (l)  ZrO2 (c)+SiO2 (l)  ZrB 2(c)+SiC(c)  C(c)  2000  Log PCO =5.0 (Pa)  Log PO2 (Pa)  g o L  P  a  r  t  i  a  l  p  r  s s e  u  r  e  (  P  a  )  Fig. 2. Volatility diagram for ZrB2-SiC at 2000 °C and PCO = 105 Pa. A stable C(c) field is present.  P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  2725          \\x0c', '2726  P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  )  a  P  (  e  r  u  s s e  r  p  r  o p  a  V  10 7 10 6 10 5 10 4 10 3 10 2 10 1 10 0 10 -1 10 -2 10 -3 10 -4 10 -5 10 -6 10 -7 10 -8 10 -9  10 -10  100 0  B2O3(g)  SiO2(g)  ZrO2(g)  A  120 0  B  1400  C  D  160 0  1800  2000  2200  2400  260 0  Temperature   Fig. 3. Vapor pressure for oxides of ZrB2-SiC. Points A and B represent the temperatures for stability and instability of B2O3, respectively. Points C and D indicate the temperatures for stability and instability of SiO2, respectively.  This may be the reason that no solid C(c) has been observed in studies of the oxidation of ZrB2-SiC in the literatures. Scale vaporization is another issue for thermodynamic stability. Note that in air (PO2 =0.2 atm), the major vapor specie above SiO2 is SiO2(g) and the minor specie is SiO(g). The major vapor species above B2O3 and ZrO2 in air are B2O3 and ZrO2, respectively. Fig. 3 shows the vapor pressure of the condensed oxides for ZrB2-SiC. ZrO2 exhibits the lowest vapor pressure of the given oxides. However, there are no guidelines for acceptable recession rates. Assuming 10\\x001 Pa as the criteria for stability, B2O3, SiO2 and ZrO2 are stable at temperatures below 1091, 1809 and 2300 °C, respectively. The temperature limits for B2O3, SiO2 and ZrO2 are 1442, 2368 and 2923 °C, respectively, assuming 102 Pa as the criteria for instability. These calculations suggest that volatility is not a major problem for ZrO2. Scale melting is also an issue for the mechanical stability and the temperature limitations for SiC and ZrB2 are 1723 and  2680 °C, respectively. This indicates that scale melting is unlikely as a major issue for monolithic ZrB2. In contrast, scale adherence is a concern for ZrB2, where the scale spallation occurs at temperatures above \\x181400 °C. From above discussion, we can conclude that the major barriers to application of monolithic SiC and ZrB2 are not the same, and in fact, most of are complementary. Therefore, the oxidation resistance can be significantly improved by merging these two materials into an optimized composition and microstructure.  3. Oxidation behaviour of ZrB2  The oxidation behaviour and mechanism of ZrB2-SiC composites are very complex especially at high temperatures. The experimental results obtained after oxidation can only describe the last stage of oxidation. However, the entire oxidation process cannot be observed from the using currently available experimental results. Therefore, the behaviour was first determined with theory and then validated experimentally. Apparently, only the oxidation behaviour of ZrB2 and SiC were understood, the oxidation mechanism of ZrB2-SiC composites can be better realized. The oxidation behaviour and mechanism of SiC have been studied extensively in the previous literatures [26,28-30], whereas that of ZrB2 is unclear. ZrB2 was found to oxidize at about 700 °C and the resulting oxide scale was dense and adhered to the bulk, thus providing an effective barrier to oxygen diffusion. Passive oxidation resistance was exhibited below about 1100 °C, while the active oxidation occurred at temperatures above about 1400 °C resulting from the rapid evaporation of B2O3 due to high vapor pressure as shown in Fig. 3. Between 1100 and 1400 °C para-linear kinetics have been observed [11]. Above 1400 °C, the rate of vaporization of B2O3 is high compared to its rate of formation. The surfaces of ZrB2 samples oxidized in the temperature range of 900-1200 °C are shown in Fig. 4. A glassy B2O3-rich phase embedding ZrO2 grains was observed (Fig. 4(a)) after oxidation at 900 °C, covering the yet not reacted material. The amount of B2O3 glass phases in the oxides after  Fig. 4. Surface SEM images of ZrB2 samples oxidized in the temperature range of 900-1200 °C. (a) 900 °C, (b) 1100 °C, (c) 1200 °C and (d) high magnification of (c). Spherical particles were also detected on the surface after oxidation of ZrB2 at 1200 °C which consisted of B and O.      \\x0c', 'P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  2727  Fig. 5. Cross-section SEM images of ZrB2 after oxidation at temperatures from 900 to 1500 °C. (a) 900 °C, (b) 1200 °C, (c) 1500 °C and (d) high magnification of (c).  oxidation at 1100 °C is lower than that after oxidation at 900 °C. A large amount of pores were found on the surface after oxidation at 1200 °C, as can be seen in Fig. 4(c). Moreover, a large number of small white spherical particles were also detected on the surface, an observation that had not been reported in previous studies. Analysis by EDS confirmed that they were mainly composed of B and O. The formation of the spherical particles is predominantly attributed to that the volatilization of B2O3 and then subsequent coagulation on the surface during the cooling process. This also supports the rapid volatilization of B2O3 at 1200 °C which is consistent with thermodynamic calculations. Interestingly, spherical particles were also found in ZrB2-SiC after oxidation at 1800 °C [22]. In this case the spherical particles were SiO2 rather than B2O3. SEM micrographs of the cross sections of ZrB2 oxidized in air from 900 to 1500 °C are shown in Fig. 5. The oxide scales formed at 900 and 1200 °C are dense and adherent to the bulk, which could impart diffusion-controlled oxidation protection to the underlying material. However, the oxide scale formed at 1500 °C shows a porous layer with open pore channels and no sign of a continuous protective layer was observed. Additionally, a crack was also detected  in the oxide scale as shown in Fig. 5(c). ZrO2 exhibits visible oriented growth at 1500 °C (see Fig. 5(d)). However, the ZrO2 grains did not sinter into integrity. The remarkable swelling of the samples observed after oxidation suggests the large volume expansion during oxidation. The large volume expansion upon conversion of ZrB2 to ZrO2 and the loss of liquid phase contribute to the formation of the crack and pores. To gain more information about the oxidation process and oxidation behaviour, higher temperature oxidation studies have been conducted. We believe that the oxidation process of ZrB2 is not only a process where ZrB2 is converted into ZrO2 and B2O3, but one that also includes migration, agglomeration and growth of ZrO2. Fig. 6 shows the clear evidence of the migration, agglomeration and growth of ZrO2. Nano-size ZrO2 particles and agglomeration were detected near the reaction interface, and did not adhere to the unaltered ZrB2. In addition, large ZrO2 particles were observed far from the reaction interface (not shown here). This means that ZrB2 was first oxidized into small particles and then was sintered into large particles after migration and agglomeration. Obviously, this phenomenon cannot be observed at low temperature because the transport of ZrO2 is negligi Fig. 6. Cross-section SEM image of (a) ZrB2 oxidized at 1800 °C and (b) the reacted zone at high magnification.  \\x0c', '2728  P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  ble and the oxide scale is adherent to the unreacted material during low temperature oxidation process. This indicates that the mobility of ZrO2 should be taken into account in the case of oxidation of ZrB2-SiC at high temperatures. Apparently, the analysis of ZrB2 oxidation is critically important to the understanding of the oxidation mechanism of the ZrB2-SiC composites.  4. Microstructure evolution of the oxide scale  Thermodynamically, both ZrB2 and SiC should oxidize when exposed to air. However, the oxidation rates of both species are negligible below about 700 °C according to previous experimental results [31]. The oxygen transport behaviour through the bulk was less well understood up to now. For dense ZrB2-SiC composites, ZrB2, SiC particles and grain boundaries are the possible routes for oxygen transport. It is commonly accepted that oxygen transport through grain boundaries is faster as compared to particles in the case of ZrB2-SiC. The presence of impurities at grain boundaries would enhance the oxygen transport. Therefore, the oxygen is first transported through the grain boundaries and then reacted with ZrB2 and SiC particles as shown schematically in Fig. 7. The oxidation behaviour and structure of the reacted layer depends significantly on the oxidation rate of these species, SiC content and distribution. When the oxidation rate of SiC is much lower than that of ZrB2 the unoxidized SiC particles become embedded in the oxide scale which can be found on ZrB2-SiC composites oxidized at temperatures below 1200 °C [31]. In the temperature range from 700 to 1200 °C, the oxide structure consisted of: (1) a B2O3 rich outer layer, (2) a subscale of ZrO2 that contained unoxidized SiC and (3) unaffected ZrB2-SiC. The main driving force for the outward growth of the B2O3 scale is the transport of B2O3 (l) from the oxidized ZrB2 to the surface due to a large volume expansion (>300%) during the transformation of ZrB2 to ZrO2 and B2O3. The liquid B2O3 formed acts as a barrier to inward oxygen diffusion. As the temperature approaches \\x181300 °C, SiC begins to markedly oxidize, resulting in the formation of a continuous surface layer above the ZrO2 layer. Meanwhile, B2O3 exhibits a rapid evaporation due to the high vapor pressure. ZrB2 and SiC have a similar oxidation rates in the temperature range of 1300 °C to approximately 1600 °C and the oxide growth is dominated by the inward movement of the reaction interface. The observed layered structure in this temperature range consisted of: (1) a SiO2 rich glassy outer layer, (2) a subscale of ZrO2 that contained some SiO2 and (3) unaffected ZrB2-SiC. The driving force for the transport of the silica is also the volume expansion upon oxidation of ZrB2-SiC which is similar to that of B2O3 at lower temperatures. Note that the mobility of ZrO2 at temperatures of 1600 °C and below is negligible which has little effect on the oxidation resistance at low temperatures. The formed ZrO2 apparently has not changed the ini O2  ZrB 2  SiC   Grain  boundaries   Fig. 8. Surface SEM image of ZrB2-20 vol.% SiC after oxidation at 1800 °C. ZrO2 with a dendritic morphology was found. The transport and growth of ZrO2 is apparent.  tial framework of ZrB2 below 1600 °C, and the structure of the subscale is similar to that of the bulk in which ZrB2 and SiC were replaced by ZrO2 and SiO2, respectively [31]. Active oxidation of SiC was observed in the temperature range of 1600-1700 °C which caused a significant increase in the oxidation rate of SiC leading to preferential oxidation of SiC in the ZrB2- SiC system. The change in the relative oxidation rates of ZrB2 and SiC alters the oxygen diffusion route, resulting in the change in the oxidation mechanism. In addition, the relative oxidation rate would change the inner structure of the oxide scale which strongly depends on SiC content. The preferential consumption of SiC particles due to active oxidation promotes the generation of pores in the oxidation reaction region. A porous layer will only develop when the amount of SiC is above the threshold for forming a 3D interconnected network. Obviously, the formation of a porous layer is detrimental to the oxidation resistance of ZrB2-SiC because the oxidation region changes from an interface to a layer accompanied by the incremental growth of the reaction area, resulting in an increased oxidation rate. The mobility and growth of ZrO2 significantly increased at temperatures of 1800 °C and above. The transport and growth of ZrO2 played an important role in the structural evolution of the oxide scale. ZrO2 with a dendritic morphology was found on the surface of the oxide scale as can be seen in Fig. 8. This is strong evidence of  Fig. 7. Schematic diagram of the oxygen transport in ZrB2-SiC composites.  Fig. 9. Cross-section SEM images of the scale for ZrB2-10 vol.% SiC after oxidation at 1800 °C. The formation of whisker layer was firstly observed.  \\x0c', 'P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  2729  Fig. 10. Cross-section SEM images of the scale for ZrB2-20 vol.% SiC with different particle size after oxidation at 1800 °C for 1 h. magnification of (a), (c) ZrB2 (2 lm)-SiC (0.5 lm), (d) high magnification of (c).  (a) ZrB2 (5 lm)-SiC (2 lm),  (b) high  the transport and growth of ZrO2 during oxidation of ZrB2-SiC. However, this feature alone cannot confirm that ZrO2 at the surface was transported from the inner oxide. ZrO2 whiskers were formed under the favourable conditions that can be differentiated easily from those of other oxides (Fig. 9). After oxidation at 1800 °C for 2 h, it was found that thickness of the surface whisker layer was obviously thicker than that after oxidation for 1 h. Apparently, the whisker layer thickness remains a constant if it was formed as result of sample’s surface oxidation. These facts further justify the transport and growth mechanism of ZrO2. As the temperature approached 1800 °C, the oxide structure changed significantly and oriented growth occurred as shown in Fig. 10. The oriented growth of the scale was primarily due to the evolution of the gaseous by-products and the flow of the silica-rich liquids, both of which promoted the growth of zirconia parallel to the direction of the gas product discharge and the flow of the liquids. It should be noted that the oxide structure is very dependent on SiC content at ultra-high temperatures and the spalling resistance of the scale formed on the medium SiC content is higher than the scale formed on the low and high SiC content. As shown in Fig. 7, SiC would be consumed first leaving pores when oxygen diffuses to the SiC particles because the oxidation rate of SiC is much faster than that of ZrB2. Thus the pores in high SiC containing composites would be interconnected where SiC particles are interconnected in three dimensions. Oxygen diffusion was dominant through the pores formed in the reaction region rather than the grain boundaries in this case because the diffusion rate of oxygen in the interconnected pores is obviously much faster than that in grain boundaries. ZrB2 grains that in contact with the pores directly exposed to oxygen and thus the oxidation simultaneously occurred from the pore boundaries to inner ZrB2 grains resulting in the shrinkage of ZrB2 grains which was in good agreement with experimental results as shown in Fig. 11. The formed ZrO2 particles are not mechanically stable under this condition and will be transport outwardly through the initially formed inter connected pores because the mobility of ZrO2 increases with increasing temperature whereas the scale adhesion decreases with increasing temperature. Thus the pore size in the inner reaction layer is larger than the original SiC particle size and would increase with increasing oxidation time, resulting in the spallation of the scale. Initial continuous ZrB2 matrix was transformed into a discrete ZrB2 structure due to partial depletion of ZrB2. In contrast, the pores would be occupied by the ZrO2 in cases where there is no short circuit diffusion path available for its transport. This is a consequence of the volume expansion that occurs upon the transformation of ZrB2 to ZrO2 leading to a decrease in the overall pores size. This explains why spallation and cracking occurred for ZrB2- 30 vol.% SiC, whereas a dense and adherent scale was formed for ZrB2-20 vol.% SiC after oxidation at 1900 °C for 1 h [23]. At high temperatures and low oxygen partial pressures, SiC would oxidize  Fig. 11. SEM image of SiC-depleted layer for ZrB2-30 vol.% SiC after oxidation at 1900 °C. Analysis by EDS confirmed that this layer was mainly composed of ZrB2.  \\x0c', '2730  P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  1.3 \\x02 103% Pa. Fig. 12. SEM cross-sectional micrographs of the ZrB2-30 vol.% SiC oxidized at 1900 °C for 10 min under the oxygen partial pressure of (a) low magnification of the oxide scale, (b)high magnification of the oxide scale. SEM micrographs and EDS analysis (not shown) indicate that the rod-like grains (marked with arrows) mainly consist of C.  according to Reaction (3) besides Reactions (1) and (2) which has not reported in the previous literature for this material system. Solid C was first observed in the ZrB2-SiC after oxidation at about 1900 °C with an oxygen partial pressure of 1.3 \\x02 103 Pa (Fig. 12). This fact is in good agreement with theoretical predictions. The dense and adherent oxide scale develops with difficulty above 2000 °C due to the high vapor pressures formed at the ZrB2/ZrO2 interface and thus the oxide skeleton (ZrO2) loses contact with the substrate leading to catastrophic oxidation of ZrB2-SiC.  5. Stability of the oxide scale and oxidation resistance  The oxidation resistance of ZrB2-SiC composites strongly depends on the oxide products formed during oxidation and a protective condensed-phase oxide scale must be stable both thermochemically and mechanically. The condensed oxide phases for ZrB2-SiC are B2O3, ZrO2 and SiO2, which contain both solid and liquid phases. The solid phase forms into a skeleton, which acts as a substrate and mechanical support to the liquid phase resulting in mechanical integrity. In contrast, the liquid phase seals the solid grain boundaries and plugs pores, leading to a restricted permeability path for oxygen. Apparently, only the coexistence of both solid and liquid phases can provide superior oxidation resistance. Pure B2O3 evaporates quickly above 1100 °C with increasing temperature. The thermochemical stability of B2O3 can be easily obtained by experiment because the temperature for its transition from stable to unstable is relatively low. However, the testing temperature required for examining the thermochemical stability of SiO2 and ZrO2 is extremely high, which is difficult to obtain in conventional laboratory furnaces and this is costly. However, it is appropriate to extrapolate the thermochemical stability of SiO2 and ZrO2 based on B2O3 according to the volatility diagrams and experimental results. B2O3 is thermochemically stable below 1100 °C and becomes unstable above 1400 °C, which corresponds to 1824 and 2300 °C for SiO2, respectively (Fig. 3). This means that the silica, once formed, can act as an effective barrier and a cohesive material at temperatures below about 1800 °C and stability will be lost at temperatures above about 2300 °C, which are in good agreement with our experimental results. There is almost no weight change for SiO2 before and after heat treatment in air below 1800 °C. However, the heat treatment of SiO2 in air at 1900 °C underwent a discernable mass loss. Additionally, no silica glass was found on the surface after arc jet testing of ZrB2-SiC at temperatures higher than 2300 °C which is consistent with thermodynamic calculations. ZrO2 exhibits significantly lower vapor pressures for all species. With the same method, we can predict the volatilization behaviour of ZrO2, which is stable below approximately 2317 °C.  The mechanical stability of the oxide scale is strongly dependent on the structure of the scale and the vapor pressure, which is largely determined by temperature and composition. The solid oxides should form a continuous skeleton where the oxide scale remains mechanically stable. In addition, the solid oxide should adhere to the base materials. This adherence is affected by the vapor pressure of BxOy. Thermodynamic calculations indicate that when the temperature approaches 2066 °C the peak equilibrium vapor pressure of B2O3 at the interface between the solid oxide and the base material could exceed 1 atm. This violently disrupts the contact of the ZrO2 skeleton with the substrate leading to spallation of the oxide scale. It should be noted that a liquid could flow, easily retreating from a region if interfacial pressures build up and this is not the case for solid. So the gaseous oxidation products for SiC with a high vapor pressure easily escape from the liquid phase and the gas pressure will not accumulate to such a high value as calculated. The significance of the gaseous products of SiC is less important than that of the ZrB2 products, although the value for the former is higher than the latter. Therefore, the volatilization of B2O3 is a major factor among these gaseous oxide products in the determining the breakdown of protection for ZrB2-SiC composites. The temperature limit for these composites is about 2066 °C based on disruption of the scale by gaseous products. However, the value for the temperature limit here refers to the temperature at the reaction interface not the surface temperature. In fact, there may be a large temperature gradient between the surface and reacted layer during aerospace applications although the thickness of the scale is relatively small. The ZrO2 is protective to at least 2300 °C according to the thermochemistry and this can rationalize why ZrB2-SiC composites exhibit good oxidation resistance after ablation at \\x182200 °C in arc jet testing. Compressive stress may be generated during the oxidation of ZrB2-SiC composites due to the transformation of ZrB2 to ZrO2, with a resulting 18.7% molar volume expansion. Scale cracking occurs when the compressive stress in the oxide layer reaches the critical failure stress of the layer. Cracking and spallation were detected at the interface between oxide scale and substrate in ZrB2-10 vol.% SiC [23]. A porous oxide scale was formed for 30 vol.% SiC-containing ZrB2 based ultra-high temperature ceramics as shown in Fig. 12. The strength and structure of the in situ formed ZrO2 skeleton is very dependent on the volume fraction of ZrB2. ZrO2 with a volume fraction in the range of 90-100% estimated from experiments helps to generate a dense and adherent scale which corresponds to 84 and 76 vol.% for ZrB2, respectively. Actual oxidation of SiC particles at high temperatures and low oxygen partial pressures occurred according to Reactions (2) and (3). This suggests that SiC does not generate the protective silica glass in the first step of the oxidation. Due to the oxygen pressure  \\x0c', 'P. Hu et al. / Corrosion Science 51 (2009) 2724-2732  2731  gradient across the scale, the SiO gas phase would condense into SiO2 where the oxygen pressure gets high enough. It is commonly believed that the consumption of SiC due to active oxidation causes the formation of pores, and that the oxidation of ZrB2 leads to the volume expansion and formation of solid porous ZrO2. These features are adverse to the oxidation resistance of monolithic SiC and ZrB2, respectively. However, in the case of the ZrB2-SiC system, these characteristics are beneficial to each other. The formation of pores in the scale as a consequence of the active oxidation of SiC could be partly compensated by the volume expansion derived from oxidation of ZrB2. Meanwhile the SiC consumption would leave behind space for the volume expansion as a result of ZrB2 oxidation leading to the decreased build up of the oxide growth stress. Moreover, the subsequent formation of silica glass will fill the residual holes, cracks and pores of ZrO2, which help to form a stable protective layer. Therefore, the oxidation resistance of ZrB2-SiC is not significantly affected by the transition from passive to active oxidation of SiC whereas it has a very deleterious effect on monolithic SiC. The SiO2 glass is retained in the ZrO2 structure due to the high wettability and strong surface tension in the porous ZrO2 skeleton [32]. The protective nature of the scale on ZrB2-SiC could be recovered by the subsequent formation of SiO2 even at extremely high temperatures after bubble bursts resulting from vapor pressure accumulation and release in the scale interior. Oxygen diffusion through the ZrO2 skeleton phase, liquid glass phase, percolating cracks, and pores are the dominant routes of the oxygen transport in the oxide scale. It is generally accepted that gas phase diffusion through the percolating cracks and pores, even in the Knudsen diffusion regime, is much easier than diffusion in condensed phases [32]. Therefore, reducing the number of defects in the oxide scale is a feasible and effective way to enhance the performance of these composites. Pore formation and size are strongly related to composition and microstructure. For a fixed volume fraction, the morphology of the porous structure is dependent on SiC particle size and distribution. Intergranular SiC grains exhibits fast oxidation rate as compared to intragranular SiC because the oxygen transport in grain boundaries is higher than that in ZrB2 grains. This in turn promotes the creation of the interconnected pores, which lead to an increased oxidation rate. The pore size on the innermost scale resulting from the preferential oxidation of SiC is related to the SiC particle size. Upon formation, the pores leave space for volume expansion resulted from the oxidation of ZrB2 which in turn causes pore shrinkage. The volume fraction and distribution of SiC can be optimized to prevent stress accumulation in the oxidized structure as well as the resulting interface debonding. Furthermore, the volume fraction and distribution of SiC can be optimized to allow easier filling pores in the oxidized structure with of silica glass and the capillary force that holds the solid and liquid phases together will be significantly increased. Decreasing pore sizes and tortuous paths also reduce the outward diffusion rate of SiO resulting in the formation of more SiO2. In addition, smaller pores and better connectivity inside the ZrO2(c) could enhance the collection and retention of the silica glass. This can explain why the smaller SiC grain sizes in ZrB2-SiC composites improves the oxidation resistance as shown in Fig. 10. This is also true for the passive oxidation of SiC at low temperatures because the finer grain size leads to a more homogeneous distribution of SiO2 in the oxide scale, which in turn decreases the inward oxygen diffusion resulting in enhanced oxidation resistance. The amount of SiC significantly affected the oxide structure above 1800 °C as suggested in Section 4. The change of the SiC content alters the oxidation mechanism, and the properties of the oxide scale. The scale adherence can be optimized by the varying the SiC content. ZrB2-SiC composites containing either a low or high content of SiC are not desireable for ultra-high temperature application because these extreme concentrations lead the break down of the protective oxide scale. For a low SiC contents the amount of formed silica glass is not enough to fill the cracks and pores of ZrO2. Moreover, the existence of compressive stress initiates the interfacial debonding between the oxide scale and base material. The high SiC content results in excessive pore formation, cracking and spalling of oxide scale. Additionally, too much SiC does not increase oxidation resistance above the melting point of SiO2 because the silica transported to the surface can easily shear off as a consequence of external force or spall in the case of bubble outbreaks. This is significantly different from lower temperature behaviour where silica glass is mechanically stable.  6. Conclusions  The oxidation mechanism of ZrB2-SiC was investigated based on theoretical analysis and experimental results. Good correspondence was obtained between both theory and experiments for the oxidation of ZrB2-SiC. The active to passive transition and the temperature limit in the oxidation of SiC was determined by a thermodynamic approach. The transition temperatures for active oxidation to form CO(g) and C(c) are 1734 and 1931 °C in air, respectively, which significantly decrease with a reduction of oxygen partial pressure. The microstructural evolution of the oxide scale and oxidation resistance of ZrB2-SiC composites were significantly affected by the temperature and composition, especially at temperatures above 1800 °C. The change in the temperature will alter the relative oxidation rates of ZrB2 and SiC, and will induce changes of the oxygen transport routes and this feature would be aggregated by the change of composition. The mobility and growth of ZrO2 increased significantly at temperatures of 1800 °C and above. SiO2 and ZrO2 are stable in air at temperatures below about 1824 and 2317 °C, respectively. SiO2 would lose its protective properties at temperatures above 2300 °C. The transport and growth of ZrO2 played an important role in the structural evolution of oxide scale. The temperature limit for the scale developed on ZrB2-SiC is significantly dependent on the vapor pressure of B2O3. The scale can be disrupted by B2O3(g) at one atmosphere partial pressure at temperatures of about 2066 °C. The strength and structure of the in situ formed ZrO2 skeleton is very dependent on the volume fraction of ZrB2 which in the range of 76-84% is helpful to generate a dense and adherent scale. For a certain fixed volume fraction, smaller grain sizes are advantageous to the oxidation properties.  Acknowledgements  This work was supported by the National Natural Science Fund for Outstanding Youths (10725207), Postdoctoral Science Foundation of Heilongjiang province, Postdoctoral Science Foundation of China and Natural Science Foundation of China (50602010).  References  [1] R. Telle, L.S. Sigl, K. Takagi, in: R. Riedel (Ed.), Handbook of Ceramic Hard Materials, Wiley-VCH, Weinheim, Germany, 2000, pp. 802-945. [2] R. Savino, M.D.S. Fumo, D. Paterna, M. 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},{
  "_id": 165,
  "PDF": "Oxidation mechanisms under water vapour conditions of ZrB2-SiC and HfB2-SiC based materials up to 2400°C.pdf",
  "Text": "['Journal of the European Ceramic Society 38 (2018) 421-432  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e l sev ie r .com / loca te / jeu rce ramsoc  Original Article  Oxidation mechanisms under water vapour conditions of ZrB2-SiC and HfB2SiC based materials up to 2400 °C  MARK  V. Guérineau⁎, A. Julian-Jankowiak  ONERA − The French Aerospace Lab, F-92322 Châtillon, France  A R T I C L E  I N F O  A B S T R A C T  Keywords:  ZrB2 HfB2 Oxidation Mechanisms  This study aims at observing and understanding the oxidation mechanisms of ZrB2-20 vol%SiC (ZS), HfB2-20 vol %SiC (HS) and HfB2-20 vol%SiC3 vol%Y2O3 (HSY) materials up to 2400 °C under water vapour conditions. After SPS sintering, fully densified samples were oxidized at several temperatures with 30 vol% H2O/70 vol% Ar during 20 s. Weight variations as well as post-test microstructural and XRD analyses allowed understanding the influence of the composition on the oxidation behavior and the evolution of each oxide sublayer. Below 1550 °C, oxidation is limited, and thin oxide layers are observed. At 1900 and 2200 °C, ZS and HS show mechanical damage (cracks, spallation), while HSY keeps its structural integrity and interlayer adherence. The addition of Y2O3 reduces the damages due to thermal stresses in the material due to the stabilization of the cubic phase of HfO2, and the formation of a Y2Si2O7 interphase that mitigates thermal expansion mismatch between the SiCdepleted layer and the HfO2 layer.  1.  Introduction  Ultra-High Temperature Ceramics (UHTCs) are good candidates for several extreme applications: thermal protection materials on hypersonic aerospace vehicle or re-usable atmospheric re-entry vehicles, specific components for propulsion, furnace elements, refractory crucibles … This family of ceramic compounds is made of borides, carbides and nitrides such as ZrB2, HfB2, ZrC, HfC, TaC or HfN which are characterized by high melting point, high hardness, chemical inertness and relatively good oxidation resistance in severe environments. Since the 2000 s studies to develop, in particular, hypersonic flight vehicles have led to a resurgence of interest for these materials. Indeed, hypersonic vehicles with sharp aerosurfaces (engine cowl inlets, wing leading edges and nosecaps) have projected needs for 2000-2400 °C materials which must operate in air and be re-usable. These conditions exceed the operating conditions for current structural materials for use in high-temperature oxidizing environments such as SiC or Si3N4-based materials, oxide ceramics and C/C composites with thermal protection, which exhibit good oxidation resistance only up to 1600 °C. Therefore, the development of structural materials for use in oxidizing and rapid heating environments at higher temperature is of great engineering importance. Moreover, UHTC materials have also high thermal conductivities, which gives them good thermal shock resistance and makes them ideal to many high temperature thermal applications. For a leading edge for example, a high thermal conductivity reduces  thermal stress within the material, by lowering the magnitude of the thermal gradient inside the part. Furthermore, it allows energy to be conducted away from the tip of the piece and re-radiated out of the surfaces of the component with lower heat fluxes. Diboride-based UHTCs also exhibit high electrical conductivity which is appreciable for manufacturing complex shape components for example (by using Electrical Discharge Machining). Many studies have proved the potential of diboride materials with 20 vol.% SiC as additives at very high temperature [1-7]. Then several authors have studied other additives to assess their influence on the oxidation resistance of the UHTC. For example, addition of TaSi2 in ZrB2-SiC compositions led to an improvement of the resistance to oxidation up to 1600 °C, but was detrimental at higher temperatures [8,9]. Addition of Y2O3 in ZrB2-SiC compositions was found to be detrimental above 1500 °C [10]. At temperatures higher than 2000 °C, authors have noted some mechanical degradations like spallation [11], non-adherence [12], or softening [13] of the oxide layer, even though sintering of the surface oxide layer might provide further protection towards oxidation at very high temperatures [14]. This paper is dedicated to the study of the oxidation mechanisms of HfB2-SiC and ZrB2-SiC based materials up to 2400 °C under water vapour conditions. After oxidation test using a home-made facility (BLOX4 [9]), post-test analyses (SEM, XRD) allow us to assess the influence of the composition, the temperature and the duration of the oxidation test on the oxidation behavior.  ⁎ Corresponding author. E-mail address: vincent.guerineau@onera.fr (V. Guérineau).  http://dx.doi.org/10.1016/j.jeurceramsoc.2017.09.015 Received 1 June 2017; Received in revised form 1 September 2017; Accepted 11 September 2017  Available online 12 September 2017 0955-2219/ © 2017 Elsevier Ltd. All rights reserved.  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Table 1 Sintering conditions, open porosity and oxidation conditions of  the studied UHTC materials.  Composition  Reference  Sintering conditions  Open porosity%  Densification rate ρ/ρth %  Oxidation conditions  ZrB2 + 20 vol% SiC ZrB2 + 20 vol% SiC ZrB2 + 20 vol% SiC ZrB2 + 20 vol% SiC HfB2 + 20 vol% SiC HfB2 + 20 vol% SiC HfB2 + 20 vol% SiC HfB2 + 20 vol% SiC HfB2 + 20 vol% SiC HfB2 + 20 vol% SiC + 3vol% Y2O3 HfB2 + 20 vol% SiC + 3vol% Y2O3 HfB2 + 20 vol% SiC + 3vol% Y2O3 HfB2 + 20 vol% SiC + 3vol% Y2O3  ZS-1200 ZS-1550 ZS-1900 ZS-2200 HS-1200 HS-1550 HS-1900 HS-2200 HS-2400 HSY-1200 HSY-1550 HSY-1900 HSY-2200  2100 °C, 5 min 2100 °C, 5 min 2100 °C, 5 min 2100 °C, 5 min 2000 °C, 5 min 2000 °C, 5 min 2000 °C, 5 min 2000 °C, 5 min 2000 °C, 5 min 1880 °C, 5 min 1880 °C, 5 min 1880 °C, 5 min 1880 °C, 5 min  0.04% 0.27% 0.09% 0.27% 0.13% 0.57% 0.15% 0.57% 10.81% 0.18% 0.13% 0.03% 0.13%  > 99 > 99 > 99 > 99 > 99 97.2 > 99 97.2 89 98.6 98.4 98.9 98.3  1200 °C-20s 1550 °C-20s 1900 °C-20s 2200 °C-20s 1200 °C-20s 1550 °C-20s 1900 °C-20s 2200 °C-20s 2400 °C-20s 1200 °C-20s 1550 °C-20s 1900 °C-20s 2200 °C-20s  2. Materials and characterizations  2.1. Materials  The following materials were selected for the oxidation tests:  (cid:129) ZrB2 + 20 vol% SiC, (cid:129) HfB2 + 20 vol% SiC, labelled ZS. labelled HS. (cid:129) HfB2 + 20 vol% SiC + 3vol% Y2O3,  labelled HSY.  Powders of ZrB2 (H.C. Starck, grade A, d50 = 2.8 μm), HfB2 (H.C. Starck, grade A, d50 = 7.6 μm), SiC (H.C. Starck, BF12, d50 = 0.6 μm) and Y2O3 (Ampere Industrie) were weighed to reach the target compositions. The powder mixtures were attrition-milled for 5 h in cyclohexane using zirconia or WC media, then dried in a rotary evaporator and sieved down to 50 μm mesh size. Sieved powders were sintered by Spark Plasma Sintering (SPS) (FCT System Gmbh, HD 125, Mateis, Lyon, France) between 1880 °C and 2100 °C (Table 1) with 7 MPa under Ar atmosphere. 2-mm-thick pellets are sintered using a 20 mm in diameter die coated with papyex. Prior to characterization tests, the graphite coating was removed.  2.2. Characterizations  The bulk density and open porosity of materials were measured by  the Archimedes method. Then, the densification level was calculated as the ratio of the apparent density on the theoretical density of the powder mixture. Values are reported in Table 1. Oxidation tests were carried out in the BLOX (Oxidation Laser Bench) facility at ONERA (Fig. 1). BLOX is a custom-made device used for oxidation tests at very high temperatures (up to 2500 °C) in controlled atmospheres (H2O, Ar, N2, H2, air…) at pressures ranging from few millibars to 4 bar [9]. Heating of samples is ensured with a highpower CO2 laser (2 kW), and surface temperature of the sample is measured with two bicolor pyrometers. Surface temperature can be monitored throughout the experiment with a dedicated software that acts on the laser power to reach the target temperature. A video camera allows the observation of the sample during testing. All the tests were carried out at a total pressure of 1 bar, under an H2O/Ar atmosphere (30/70 vol%, in Standard Liter). First, vacuum is made inside the chamber. Then, argon is introduced and water vapour is injected in the chamber through a peristaltic pump up to 1 bar, and the chamber itself is thermostated at 150 °C to keep water gaseous. Thus, the only oxidizing species in the atmosphere is water vapour. Pyrometers can measure surface temperatures from 1000 to 2500 °C. Once 1000 °C is reached (laser power ramp), a temperature ramp of 5 °C/s is imposed, allowing the sample to reach the target temperature (1200, 1550, 1900, 2200 or 2400 °C). After a dwell time of 20 s, the temperature is decreased at the same rate. Examples of temperature profiles and laser power profiles vs time are presented in Fig. 1. During  Fig. 1. (left) Schematic view of the BLOX apparatus, (right) Surface temperature and laser power vs time during oxidation of HS-1900, HSY-1900 and ZS-1900.  422  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 2. SEM micrographs of a) ZS-1200, b) HS-1200 and c) HSY-1200. The different layers are 1) ZrO2 + SiC, 2) HfO2 + SiC and 3) bulk material.  Fig. 3. SEM micrographs of a) ZS-1550, b) HS-1550 and c) HSY-1550. The different layers are 1) SiO2-rich glassy layer, 2) ZrO2 + SiO2, 3) HfO2 + SiO2 and 4) bulk material.  the laser power ramp (under 1000 °C), the laser power is limited at 350 W. This explains the laser power plateau observed for the HS-1900 sample, between 140 and 210 s. Samples are weighed before and after tests to measure mass variations and degradation rate. After tests, they are impregnated in an in half, polished down to 3 μm and a carbon layer is epoxy resin, cut deposited on the surface before SEM observations. Identification of the nature and thicknesses of the layers of the oxidized layers was performed using SEM (DSM 982, Zeiss) equipped with EDS. Phases on the surface were characterized using an Empyrean θ-θ). The Panalytical diffractometer (Bragg-Brentano configuration, diffraction range was from 8° to 94° with a 0.01° scanning step. Semiquantitative analyses were performed through Rietveld refinement using MAUD® software.  3. Results  First, the interest of short dwell-time (20s) is the observation of early phenomenon that could be smoothened with time. As it will be further detailed, microstructure of the oxidized layers is likely dependent of two main parameters: test temperature and the nature of the samples. At 1200 °C, SiC is not oxidized in any tested sample, except HSY1200 where SiO2 is scarcely detected at the surface. HfO2 is scarcely  detected on the top surface of HS but XRD results show principally HfB2. A thin layer of monoclinic HfO2 (5 μm) or ZrO2 (5 μm) is detected respectively on HSY and ZS materials (Fig. 2). At this temperature, a very thin oxidized layer is developed. The oxidized surface retained the microstructure of the MeB2-SiC material. No glassy layer is observed on the surface. At 1550 °C, each material exhibits a two-layered oxide scale: on the very top is a glassy SiO2-rich layer, and beneath this layer is a MeO2 (Me = Zr or Hf) layer (Fig. 3). At this temperature, no mechanical issue is observed, neither any SiC-depleted layer (i.e. a MeB2 layer without SiC). At 1900 °C, each material exhibits a three-layered oxide scale: from surface to bottom, a heterogeneously distributed SiO2-rich glassy layer, a MeO2 layer, and then a SiC-depleted MeB2 layer. However, some microstructural differences can be noticed. On the ZS sample, The ZrO2 layer shows two different aspects. At the very top, non-cohesive lamellae are observed on the edges of the oxide scale. Due to its non-cohesive nature, this sublayer fell off the sample during handling, and was consequently not observed in the SEM micrographs. It is nonetheless assumed that it was present during the oxidation test. Beneath that non-cohesive ZrO2 layer, lies a ZrO2 layer, made of large ZrO2 grains and columnar ZrO2 infiltrated by SiO2-rich glass. Below this last ZrO2 layer is the SiC-depleted ZrB2 layer, which is on top of the unreacted material (Fig. 4 and 5). Si-O-C rich inclusions are found in the SiC-depleted layers, along with empty pores.  Fig. 4. SEM micrographs of: a) ZS sample oxidized for 20 s at 1900 °C. Zoomed-in micrographs of the same sample are labelled a-1) and a-2). Different layers are: (1) ZrO2 lamellae, (2) ZrO2, (3) SiC-depleted ZrB2 and (4) bulk material. Silica-rich glass is observed in layers (1) and (2). b) Zoom on big-sized ZrO2 grains in layer 2.  423  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 5. SEM micrograph showing the SiC-depleted layer of  the ZS-1900 sample. Different  layers are (1) ZrO2 infiltrated by silica-rich glass, (2) SiC-depleted ZrB2 where empty pores and Si-O-C inclusions are found, and (3) bulk ZrB2-SiC material.  Microstructure of the oxidized HS sample is quite different (Fig. 6) and the oxide scale is irregular. One can observe the presence of “volcano-like” structures, with SiO2-rich glass erupting from the curved HfO2 layer. Under those structures, a sublayer of HfO2 is present. The sample is thus divided in two distinct zones: the flat zones where the 20 μm thick HfO2 layer lies right above the SiC-depleted HfB2 layer, the 20 μm thick HfO2 lies and the volcano-like zones, where under another sublayer of HfO2, itself being above the SiC-depleted HfB2 layer. Many cracks are also observed within the oxide scale (Fig. 6b) On  the contrary, oxidized HSY sample exhibit a regular oxide scale, with no apparent decohesion. The scale consists on a top SiO2 glassy layer, then a columnar HfO2 layer, above a SiC-depleted HfB2 layer (Fig. 6c). Si-OC and Y2Si2O7 inclusions are found in the SiC-depleted layer (Fig. 6d). At 2200 °C, the top oxide scale developed by oxidized ZS sample also shows two different aspects. Same non-cohesive lamellae as ZS1900 are observed at the very top but fell off the sample during handling. Beneath that non-cohesive ZrO2 layer lies a columnar ZrO2 layer, then the SiC-depleted ZrB2 layer and the unreacted material  Fig. 6. SEM micrographs of (a, b) HS-1900 and (c,d) HSY-1900. The different layers for HS-1900 are: (1) irregular SiO2-rich glassy layer, (2) HfO2, (3) SiC-depleted HfB2 and (4) bulk material. For HSY-1900, the different layers are: (5) SiO2-rich glass, (6) columnar HfO2 filled with silica, (7) SiC-depleted HfB2 with black Si-O-C rich inclusions and (8) bulk material. White arrows in (b) point at cracks. In (d), inclusions of Si-O-C (black) and Y2Si2O7 (gray) are shown inside the SiC-depleted layer.  424  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 7. SEM micrographs of ZS-2200. Different layers are: (1) ZrO2 lamellae, (2) columnar ZrO2, (3) SiC-depleted ZrB2 and (4) bulk material. Silica-rich glass is observed in layers (1) and (2).  the columnar  (Fig. 7). Compared to the material oxidized at 1900 °C, ZrO2 layer is infiltrated with SiO2. Contrary to the microstructure observed at 1900 °C, the oxide scale developed by the HS sample at 2200 °C is more regular: it consists on a top porous HfO2, non-adherent to the subjacent SiC-depleted HfB2 layer. Inside the pores, some silica-rich glass with HfO2 inclusions are observed. Porosity on the top of the SiC-depleted layer is filled with silica-rich glass (30 μm). C-rich inclusions are also found deep within this layer (Fig. 8). The oxide scale developed by the HSY sample consists of a top HfO2 layer, above an HfB2 layer which porosity is filled with a Y-O-Si-particles rich glass, identified as Y2Si2O7 by EDS. Beneath this layer is the SiC-depleted HfB2 layer. Interestingly, some spallation is observed, but is located within the top HfO2 layer, which is cut in two parts (Fig. 9). Only HS was oxidized at 2400 °C. Oxidized sample developed a top monoclinic HfO2 with no or very little porosity, non-adherent to the subjacent SiC-depleted HfB2 layer (with Hf-O-C inclusions) (Fig. 10). This suggests that, through intense heating, the top layer has sintered  during the test. Samples surfaces were analyzed via XRD, in order to determine the dominant crystalline structure of the −if presentoxide layer (Fig. 11a, b and c). Concerning ZS sample, monoclinic ZrO2 was detected at 1200 °C, coexisting with ZrB2 (roughly 80/20 wt%). With increasing temperature, monoclinic ZrO2 became more and more preponderant. Concerning HS sample, HfO2 signal was not detected at 1200 °C. Actually, the XRD diagram of an HS sample oxidized at 1200 °C was very similar to the unoxidized HS sample. At 1550 °C, monoclinic HfO2 is detected, along with HfB2 (roughly 93/7 wt%). The same trend is then observed when temperature is increased, and only monoclinic HfO2 is detected at 2200 °C and 2400 °C. While ZS and HS developed an only monoclinic MeO2 layer after oxidation, this was not the case for HSY. From 1550-2200 °C, both monoclinic and cubic HfO2 can be seen on the X-ray diagrams (no tetragonal hafnia was observed on any diffractogram). Moreover, the addition of Y2O3 had an impact on the crystalline structure of the oxide layer during heating, and consequently on the microstructure of the oxidized samples. First, Non-oxidized HSY  Fig. 8. SEM micrographs of HS-2200. Different layers are: (1) HfO2, (2) SiC-depleted HfB2, (3) bulk material. detected across the SiC-depleted HfB2-layer.  In the HfO2 top layer, silica rich glass areas (4) are detected. Cracks are  425  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 9. SEM micrographs of HSY-2200. Different layers are: (1) HfO2, (2) SiC-depleted HfB2 filled with Y-Si-O glass and (3) SiC-depleted HfB2. a-1), a-2) and a-3) are the corresponding zoomed-in micrographs of a).  sample and HSY-1200 had a peak at 2θ = 26.5°: this peak can either be attributed to graphite or quartz. But HSY-1200 also had a peak at 2θ= 20.2° (Fig. 11g), which is only attributed to quartz. The graphite XRD peak can be attributed to papyex from SPS molds. Even if the surface is cleaned before test, some residues can still be present. HfSiO4 is detected for HSY-1550, and results from the reaction between SiO2 and HfO2 in the external layer. As HfSiO4 melts incongruently at 1760 °C, it is not detected for the upper temperatures. At 1900 °C, we must take into account that surface analysis will mostly give information on the very top layer. Many inclusions are detected in the external glassy layer of HSY-1900 (Fig. 12) and EDS analysis revealed the presence of Hf, O and Y. Thus, these inclusions are most likely yttria-stabilized hafnia, which explains the large majority of fluorite HfO2, compared to monoclinic HfO2, for this sample. The slight shift of the main fluorite peak towards the lower angles can be explained either by an higher doping of yttria in the hafnia structure, or by the presence of Y2Hf2O7 phase, which main peak overlaps with the fluorite. But as the other peaks of Y2Hf2O7 are not observed, the yttria doping is the most likely explanation. The presence of fluorite HfO2 at the surface of the HSY-2200 sample demonstrates that cubic HfO2 is also present in the  bulk HfO2 layer, as this sample does not show any surface glassy layer. The broad diffusion peak of a SiO2-rich glassy phase is visible on the diffractogram of each sample oxidized at 1900 °C (Fig. 11d,e,f). This is consistent with the SEM micrographs of these samples, in which a glassy layer is observed at the surface (Fig. 4 and 6). Moreover, HSY1900 exhibits peaks corresponding to orthorhombic Y2Si2O7. This phase is not detected at lower or higher temperature.  4. Discussion  The predominant oxidizing molecules is water vapour. Thus, expected oxidation mechanisms under at atmospheric pressure are as follow: [15-17]  ZrB2(s) + 5H2O(g) = ZrO2(s) + B2O3(l) + 5 H2(g)  HfB2(s) + 5H2O(g) = ZrO2(s) + B2O3(l) + 5 H2(g)  SiC (s) + 3H2O(g) = SiO2(l) + CO(g) + 3 H2(g)  Vaporization reactions:  SiO2 (l) = SiO2 (g)  (1)  (2)  (3)  (4)  Fig. 10. HS sample oxidized at 2400 °C for 20 s. Different  layers are:  (1) monoclinic HfO2 and (2) SiC-depleted layer. Cracks are present across the SiC-depleted HfB2 layer.  426  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 11. X-ray diffractograms of a) ZS, b) HS and c) HSY samples oxidized at different temperatures. On d), e), f) broad diffusion peaks of SiO2 rich glassy phase of samples oxidized at 1900 °C. On g), the SiO2 (quartz) peak present on HSY-1200, and absent on unoxidized HSY.  Fig. 12. SEM micrographs of the silica-rich layers of  (a) HS-1900 and (b) HSY-1900.  B2O3 (l) = B2O3 (g)  and the most likely vaporization reactions with water vapour:  B2O3(l) + H2O(g) = 2 HBO2(g)  SiO2(l) + 2H2O(g) = Si(OH)4(g)  (5)  (6)  (7)  Due to the short duration of the tests, no significant oxide layers developed at 1200 °C (Fig. 2). At this temperature, SiC was not oxidized. Interestingly, no glassy layer was observed at those temperatures. It is known that water vapour environment favors volatilization of boron-oxide compounds, for instance HBO2(g) following Eq. (6) [18]. However, it must be noted that EDS analysis of the HSY-sample showed the presence of some conglomerates of SiO2 on its surface, likely quartz, according to the XRD analysis.  4.1.  Influence of  the material on oxidation behavior  At 1550 °C, each oxidized sample exhibits a thin silica-rich glassy layer, as SiC is oxidized at those temperatures. Even though water vapour favors volatilization of silica and the oxidation rate of SiC [19] at high temperatures, a glassy layer is nonetheless formed (Fig. 3). An interesting feature of the very top of the ZrO2 layer of ZS-1550 is that it is made of fine-grained zirconia (Fig. 13, a). According to Karlsdottir  and Halloran [20,21], these fine-grained zirconia may be formed by precipitation of the zirconia in the BSZ (boria-silica-zirconia) liquid due to the evaporation of boria during the test. As B2O3 evaporates, the BSZ liquid moves into the two-phase region with solid ZrO2 and BSZ liquid. Then, the precipitated zirconia gets deposited on the top of the oxide layer, thus forming an outer layer of fine-grained zirconia. This mechanism is plausible in this case, as we know that water vapour enhances the vaporization of B2O3(l). Inclusions are also observed in the glassy layer HSY-1550 (Fig. 13b). In this case, EDS showed that these inclusions are mainly HfO2 with a small quantity of yttria. Thus, and in agreement with the XRD diffractograms, those inclusions are very likely crystallites of cubic hafnia. At temperatures higher than 1900 °C, ZS, HS and HSY behaved very differently. The top ZrO2 layer is divided into lamellae, which were so fragile that it fell off the sample during handling. However, these lamellae can be seen on the edges of the sample (Fig. 4). Beneath those lamellae, a ZrO2 layer made of columnar zirconia and big-sized grains, the growth of grains being favored by the high oxidation temperature, as observed for instance by Han et al. [14]. This layer is, for the most part, non-adherent to the subjacent SiC-depleted ZrB2 layer. Glassy SiO2 is present on the top of the lamellae, as well as on the columnar ZrO2 layer and the XRD diagram shows an obvious broad diffusion peak corresponding to silica-rich glassy phase (Fig. 11f), but it does not form a surface layer, as can be seen for HS and HSY. Many dendritic  427  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 13. SEM micrographs of (a) the top layers of ZS-1550 and (b) the silica-rich layer of HSY-1550. Different layers are: (1) SiO2-rich glassy layer and (2) ZrO2 infiltrated by silica-rich glass.  also present in the bulk hafnia layer. The presence of cubic HfO2 at low temperatures could decrease the effect of the volume contraction caused by the monoclinic to tetragonal transformation. This would explain why there are no internal cracks or spallation observed in the oxide layer of HSY sample oxidized at 1900 °C, while all these mechanical flaws are observed for the HS sample oxidized in the same conditions. At this temperature, oxidized HSY sample managed to keep its structural integrity, with a top HfO2 layer presenting a columnar structure that allows easier evacuation of the high-pressure gases and the glassy borosilicate, which are oxidation products. Under the same conditions, oxidized HS sample did not keep its structural integrity, oxidation products being likely evacuated preferentially through the volcano-like structures and the internal network of cracks, the latest being filled with SiO2-rich glass. In addition, Li et al. [23] who worked on the oxidation of ZS samples, showed that when a columnar ZrO2 framework was infiltrated by a borosilicate phase, its mechanical resistance was improved compared to a bulk material thanks to the fact that the glassy phase could mitigate the pressure build-up by regulating the dynamics of gas bubbles. In this case, the HfO2 framework developed by HSY is infiltrated by the silica-rich glass, and this effect can also explain its better mechanical integrity. Glassy outermost layers developed by HSY and HS samples oxidized at 1900 °C also exhibit some differences: as described earlier, HSY exhibits an homogeneous and flat SiO2-rich glassy layer, which is continuously replenished by borosilicate glass coming through the columnar pores of the HfO2 layer, while HS exhibits an heterogeneous glassy layer, this heterogeneity being due to the fact that the borosilicate comes outwards preferentially through the volcano-like structures. The density of inclusions inside the glassy layer is also a major difference. As can be seen on Fig. 12, glassy layer on top of the oxidized HSY sample contains a lot of inclusions, which according to EDS measurements are Hf-Y-O inclusions (mostly Hf and O). Interestingly, EDS measurements pointed on the crystallites inside the HSY1900 glassy layer always showed the presence of a small amount of yttrium, while yttrium was absent of the hafnia situated on the hafnia layer beneath the glassy layer. The borosilicate can dissolve both HfO2 and Y2O3. Upon vaporization of the borosilicate layer, or upon cooling, crystallization will occur, in the form of yttria-stabilized hafnia, which structure is fluorite. Thus, difference must be made between “primary” hafnia that results from the oxidation of HfB2, which is mostly monoclinic at room temperature, and “secondary” hafnia, that results from the recrystallization, which is mostly under its cubic form. This is attested by the XRD analysis (Fig. 11c), where it can be seen that, for HSY-1900, the fluorite peaks are clearly predominant, while this is not  Fig. 14. SEM micrograph showing dendritic structure in the oxide layer of ZS-1900.  structures are evidence for recrystallization of ZrO2 (Fig. 14). The oxide scale developed by the HS sample shows some convection patterns. Even though the HfO2 layer is, for the most part, adherent to the subjacent SiC-depleted HfB2 layer, there are punctually some protuberances consisting of the HfO2 layer being non adherent to the underlayer (Fig. 6). It can be seen that such structures are a direct outwards pathway for the glassy SiO2, and is characteristic of a convection mechanism. But it is also an inward pathway for oxidizing molecules, as evidenced by the HfO2 layer which is beneath the protuberance. This HS sample also exhibits a large amount of fractures (Fig. 6). Even though water vapour environment is known to favor volatilization of the silica-rich glassy layer, it is likely replenished during oxidation, as evidenced by the huge thickness of the SiC-depleted layer (92 μm, to be layer which is 28 μm in compared with the thickness of the HfO2 average), layer in which SiC is probably oxidized to form SiO(g), which then condensates near the surface into SiO2. Contrary to ZS and HS, the oxide scale developed by HSY shows no mechanical defaults and seems more homogeneous. The 3 vol% Y2O3 addition (compared to HS) is therefore a key parameter in the final microstructure. Stacy studied the yttria-hafnia system with X-ray diffraction [22], and showed that when a Y2O3-HfO2 (1/99 mol%) was annealed for 4 h at 1600 °C, both cubic and monoclinic would coexist. This is in agreement with our XRD characterizations (Fig. 11, c). A factor that must be considered for the cracks observed in the oxide layer of HS sample oxidized at 1900 °C is the volume contraction due to the monoclinic to tetragonal configuration. This transformation occurs around 1700 °C [22]. But analysis of the XRD diagram of the HSY sample oxidized at 1550 °C showed that there was already more than 30 wt% of cubic HfO2. Even though it is believed that most of the fluorite hafnia is situated in the outer glassy layer, we assume that it is  428  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 15. SEM micrographs of different aspects of the ZS-2200 sample.  the case for HSY-2200, because there is no glassy layer, which appears to be the main medium for the growth of cubic yttria-stabilized hafnia. Cubic hafnia which is present at the surface of HSY-2200 may come from the redeposition on the surface of the crystallites once the glassy layer is completely evaporated. This explains why the cubic hafnia appears at 1550 °C, while it is absent at 1200 °C: the first exhibits a glassy layer with many crystallites, when the second does not present any glassy layer on its surface. This also explains why HSY samples have cubic hafnia on their surface, when HS samples do not. Inside the SiCdepleted HfB2 layer, SEM/EDS characterizations showed the existence of Y-O-Si rich inclusions that were identified with EDS as Y2Si2O7. XRD diagrams on the surface of this sample showed evidence of the presence of orthorhombic Y2Si2O7, a high-temperature phase of this compound. The SiO2-Y2O3 phase diagram shows the existence of an eutectic phase at 1645 °C. This Y2O3 rich borosilicate glass then flows outwards, where B2O3 and SiO2 preferentially evaporate, thus favoring precipitation of Hf-Y-O rich inclusions. At 2200 °C, ZS-oxidized sample exhibits a similar microstructure (with thicker layers) than the ZS sample oxidized at 1900 °C. As described earlier, the top layer of zirconia spalled off the sample, and therefore could not be seen on SEM observations. However, the formation of this spalled layer can be observed thanks to the videocamera pointing at the sample during oxidation. When the sample reaches 1900 °C, a hot spot suddenly appears on the sample. This hot spot is located where the spalled zirconia layer was, before falling off the sample. At this temperature, the pressure buildup of volatile oxidation products is high enough to spall this layer off. The same phenomenon was observed for ZS-1900, and explains why, in the surface temperature vs time profile (Fig. 1), we observed a 50 °C overshoot when the sample reaches the 1900 °C target temperature: the spallation of the zirconia suddenly created an insulating layer where the pyrometer was pointing at, and the temperature regulation system did not react quick enough to this surface property modification. The formation mechanism of zirconia lamellae is unsure. Considering the high porosity of  the zirconia layer, the solubility of zirconia in the silica-rich glass must be very high in this water-vapour atmosphere. Han et al. studied the oxidation of a ZrB2-20 vol%SiC, using an oxyacetylene torch, at 2200 °C for 10 min [14]. They reported the formation of an outer layer of zirconia made of large particles, assuming that this dense coherent layer could enhance oxidation resistance. This is obviously not the case in our case. As can be seen on Fig. 15a), longitudinal porosity is observed inside the external zirconia layer. As silica evaporates, the upper layer form “macro-lamellae”. This can spall off the lower layer and thus spallation if favored by the high pressure of the gases formed during oxidation such as SiO(g), HBO2(g), SiO(OH)4(g) or H2(g) (Eqs. (1-7)). The formation of this longitudinal porosity might be due to the infiltration of the glassy layer in the intergranular porosity between zirconia grains (Fig. 15b). With further dissolution of zirconia in the glassy layer, this filled porosity expands, and if the silica-rich layer evaporates completely, then two distinct layers of zirconia are formed. As can be seen on Fig. 7, some lamellae are very thin. Their formation might be due to the precipitation of secondary zirconia, precipitation being triggered by the continuous evaporation of the silica-rich layer. A fine lamellae in formation is shown in Fig. 15c). Those lamellae are very fragile, and these structures are not seen on the center of the sample (where the temperature was the highest). The HS sample exhibits an oxide layer that can be roughly divided into two parts: a top monoclinic HfO2 layer that is not adherent to the rest of the material, and a SiC-depleted HfB2 layer. The top HfO2 does not seem columnar and presents a lot of pores with a very large size range (few μm to hundred μm large). Open pores are empty, while closed ones retained SiO2-rich glass. It is likely that borosilicate evaporated through the open porosity. When glass was retained in the closed pores, a lot of recrystallization occurred, thus explaining the presence of many inclusions in the SiO2-rich glass, along with many dendritic features (Fig. 16a). In the oxidized HS sample, the decohesion was clearly seen between HfO2 layer and HfB2 layer. This is not the case in the oxidized HSY  429  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  Fig. 16. SEM micrographs of (a) the top HfO2 layer of HS-2200 and (b) HSY-2200. (a) Arrow points at dendritic features, evidence of recrystallization of HfO2 in the silica-rich glass. (b) The different layers are: 1) HfO2, 2) HfB2 + Y2Si2O7 and 3) SiC-depleted HfB2 layer.  sample: in this case, the decohesion is located within the HfO2 layer, and does not extend to the whole sample surface. Under the HfO2 layer is an HfB2 layer, whose porosity is filled with a Y2Si2O7-rich glass (Fig. 16b).  4.2. Analysis of oxidation features  From 1900 °C, each sample developed a SiC-depleted layer in the oxide layer, i.e. a layer where SiC is absent, but the MeB2 matrix remained unchanged. According to the nature of the sample, the characteristics of this layer changed. First, the size of this layer, compared to the whole oxide layer, is very dependent on the sample nature. For oxidized HS samples, the SiC-depleted layer constitutes 70-80% of the 50-60% and oxide layer. For HSY, for ZS, it falls down to 20% (Table 3). Several tendencies can be observed through analysis of the thicknesses of oxide scales and sublayers (Table 3). Oxidized ZS samples have overall smaller oxide scales for a given temperature, and ZS samples have very thin SiC-depleted layers, compared to HS and HSY. If we compare the MeO2 layers, HS samples show the smallest thicknesses. When ZS is oxidized under water vapour atmosphere at temperatures higher than 1900 °C, the top ZrO2 layer can be divided in two parts: a very top layer, that mostly spalled off the sample (after the test), and the bottom layer, which is made of columnar ZrO2 and big-sized zirconia grains. Zirconia is an insulating material, and it is likely that the top ZrO2 acted as a thermal barrier. The insulating properties must have been enhanced by the fact that it was non adherent thus leaving a  Table 2 Weight variations of oxidized samples in function of the temperature.  1200 °C  1550 °C  1900 °C  2200 °C  2400 °C  ZS HS HSY  0.04% 0.01% 0.01%  0.15% 0.01% 2.24%  0.00% −0.08% −0.28%  −0.32% −1.27% −1.48%  −1.57%  layer of atmosphere between the top ZrO2 layer and the columnar ZrO2. Moreover, ZrB2 has a smaller thermal conductivity compared to HfB2 (respectively 60 and 104 Wm−1.K−1, [24]), so the in-depth temperature might be smaller for ZS samples compared to HS samples, which might explain why the SiC-depleted layers of ZS samples are much smaller (Table 3). Another argument is that, at these high temperatures, diffusion of oxidizing molecules through the ZrO2 layer cannot be neglected (while diffusion through HfO2 is much lower [25]). But for a SiC-depleted layer to form, the partial pressure of oxidizing molecules must be low enough to promote active oxidation of SiC without oxidizing ZrB2 [26]. So if the overall permeation of oxidizing molecules is high enough, then the formation of a SiC-depleted layer is not promoted. This explains why, compared to oxidized HS and HSY samples, oxidized ZS sample has such relatively low SiC-depleted layers. Those results are overall confirmed by the weight gains/loss analysis (Table 2), where it can be seen that ZS samples lose less weight during oxidation tests. If weight loss can be interpreted, among other factors, as the departure of SiC through active oxidation, then the very small SiC-depleted zone and the overall smaller oxide scale (Table 3) explains the weaker mass loss. Plus, assuming that the borosilicate is fueled by SiO(g) formed by active oxidation of SiC which then is reoxidized into SiO2(g) [27], then the small thickness of the SiC-depleted, compared to the total oxide scale, might lead to a borosilicate glass that is too rich in B2O3, which is unstable at high temperatures (and mostly under water vapour environment) and less viscous than a richer-in-SiO2 borosilicate, thus increasing the diffusivity of oxidizing molecules through this layer. Comparing HS and HSY-oxidized samples, one can see that HSY samples exhibit thicker top HfO2 layers than HS samples. At 1900 °C, for HSY sample, the top HfO2 layer exhibits a columnar structure, filled and covered with SiO2-rich glass, while for HS sample, the top HfO2 layer exhibits a denser microstructure with some scarcer porosity. As assumed by Carney [5], the columnar growth during the oxidation of HS and ZS samples is triggered by the flow of MeO2-SiO2-B2O3 glass and its convection. As it can be seen on Fig. 12, a lot of recrystallization occurred during cooling in the silica-rich layer of HSY-1900, but not in HS-1900. Therefore, the glass formed during oxidation of HSY  Table 3 Thicknesses of the total oxide layer (SiC-depleted layer included) and thicknesses of the SiC-depleted layer for different samples at different oxidation temperatures. In brackets is the ratio between the SiC-depleted layer and the total oxide layer. At least ten measurements were made for each value.  1900 °C  Total oxide layer (μm)  SiC-depleted layer (μm)  Silica-rich glassy (layer)  Total oxide layer (μm)  SiC-depleted layer (μm)  2200 °C  2400 °C  Total oxide layer (μm)  SiC-depleted layer (μm)  ZS HS HSY  114a ± 23 120-150 205 ± 12  26 ± 9(23%) 92 ± 13(60-77%) 121 ± 6(60%)  - 0-30b 12 ± 3  291* ± 36 546 ± 13 515 ± 24  51 ± 16 (18%) 395 ± 7(72%) 270 ± 23(52%)  N/A 897 ± 75 N/A  N/A 676 ± 53(75%) N/A  a  : The thickness of these layers were measured slightly off the center of the sample, where the oxide layer kept its integrity. b : Due to the heterogeneity of the silica-rich layer, only the extreme values are mentioned.  430  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  solubilized more HfO2 than in HS samples. This explains why HSY-1900 exhibits a columnar HfO2 layer, when HS-1900 does not. The columnar structure is a better pathway for the oxidizing molecules, compared to a non-oriented porosity. Moreover, Tian et al. [28] showed that small additions of Y2O3 to a borosilicate glass would lower its viscosity at high temperatures. Similarly, Shimizu et al. have shown that Y2O3 additions in a binary MgO-SiO2 glass would lower its viscosity [29]. It is therefore possible that dissolved yttria would lower the viscosity of the borosilicate glass in the case of the oxidized HSY sample. The great amount of Y-rich particles in the glass phase (Fig. 12) that crystallized during cooling or by the glassy layer evaporation show that the glassy phase solubilized HfO2 and Y2O3 during oxidation. A lower glass viscosity would in turn increase the permeability of this phase towards oxidizing molecules. Thus, the partial pressure of oxidizing molecules beneath the top HfO2 layer may be greater for HSY sample than HS sample. If this partial pressure is high enough, then oxidation of HfB2 is favored, and thicker HfO2 layers are consequently observed. This interpretation is assessed by the convection patterns observed on the top of the oxidized HS-1900 sample: underneath the protuberance in which glass and oxidizing molecules can easily get through, a layer of HfO2 is observed, meaning that if the passage of oxidizing molecules is favored, then the oxidation of HfB2 is favored, rather than only active SiC oxidation. Beneath those convection patterns, the SiC-depleted layer is smaller compared to the ones developed under the “flat” zones of the same oxidized samples. To summarize, due to their very different microstructures, the top HfO2 layer developed by HSY is thicker and more permeable to oxidizing molecules than the HfO2 layer developed by HS, mostly because of the columnar porosity of the top HfO2 layer of HSY-1900. In the end, once the top HfO2 layers developed by HSY and HS are formed, they provide a barrier towards oxidizing molecules, so that the partial pressures of oxidizing molecules beneath is low enough to promote the formation of thick SiC-depleted layers. At 2200 °C, none of these oxidized samples exhibit an external glassy layer. Likely, the silica-rich glass evaporated due to the high temperature of the water vapour-rich atmosphere. The microstructures of their top HfO2 layer show some similarities: none of them exhibit columnar structure, but rather a slightly porous homogeneous layer, with great disparities in the pore sizes. However, strong differences exist: the top HfO2 layer developed by the HS sample remained coherent, but spalled off the underneath SiC-depleted layer, while the top HfO2 layer developed by the HSY sample remained adherent to the layer underneath, but is on a great portion divided in two. Plus, concerning HSY, the layer directly under the top HfO2 layer is not the SiCdepleted layer (where SiC is “replaced” by empty pores), but an HfB2 layer which porosity is filled with Y2Si2O7 glass. According to Courcot et al. [30], who performed thermodynamic study of the Y2O3-SiO2 system, yttrium silicates are more stable than the single oxides at high temperature under moist environment. This might explain why Y2Si2O7 was observed in the oxidized sample, instead of a SiO2rich glass. This layer might be responsible for the improved adherence between the HfO2-rich and the HfB2-rich layers. Carney et al. [31] studied the oxidation of HfB2-SiC and HfB2-MoSi2 at 1950 °C. The first material exhibited a spallation of the oxide scale, while the second material exhibited an adherent oxide scale. For the second material, a Mo-B phase was observed at the interface, and this phase could have led to a better match between the CTE of the bulk and oxide scale, thus improving adherence. According to the literature, the average CTE of: HfB2 is 6.3 10−6/°C [32], monoclinic HfO2 is 7.46 10−6/°C (at 1200 °C (α11 = 8.86, α22 = 1.74 [31]) but with strong anisotropy and α33 = 11.77 10−6/°C), cubic yttria-stabilized HfO2 is 10.5 10−6/°C [33] (for HfO2/Y2O3 90/10 wt%, from 24 to 1482 °C). In our case, the Y2Si2O7 embedded in an HfB2 layer is situated between the HfO2 layer (monoclinic and cubic) and the SiC-depleted layer. The CTE of orthorhombic δ-Y2Si2O7 is 8.1 10−6/°C [34]. With a mixture law, and assuming Y2Si2O7 occupies 20 vol% of the layer, the  average CTE of the HfB2-Y2Si2O7 layer would be 6.7 10−6/°C. Thus, the CTE of this layer is slightly higher than the one of the underneath SiCdepleted HfB2 layer ( 6.3 10−6/°C) and lower than the above HfO2 layer ( > 7.5 10−6/°C). This explains why there is no mismatch between the SiC-depleted layer and both of these layers. By creating an interface with intermediary CTE between the SiC-depleted HfB2 layer and the HfO2 layer, it may avoid decohesion in the oxidized HSY sample while this decohesion is observed in oxidized HS samples, where there is no interface between the SiC-depleted layer and the HfO2 layer. Concerning the huge difference in size of the top HfO2 layer, this must be due once again to different overall permeability to oxidizing molecules. On the HSY sample, one can see that there are two different structures for the top HfO2 layer: at the center of the sample, the HfO2 is split in two parts, with small-sized pores, while on the edge, the HfO2 layer is not split, but exhibits big-sized pores. Based on this observation, the formation of the two-parts HfO2 layer might be due to the coalescence of the pores, forming one big flat pore which then acts as a privilieged pathway for oxidizing molecules. At the very center of the sample, the top layer is crossed by a pore, which also acts as a privilieged pathway, as was seen on the HS sample oxidized at 1900 °C. At 2400 °C, one can see that the HfO2 layer has almost no porosity on the top, and much more on the bottom. Those two different sublayers are separated by a gap. Due to very high temperature, HfO2 have sintered, thus leaving no or very few porosity on the outermost surface. Such sintering of the top oxide layer has been observed by others [14]. The crack that separates the layer could be due to the CTE mismatch between the non-porous HfO2 and the porous HfO2. Transversal cracks are seen across the SiC-depleted HfB2 layer.  5. Conclusion  ZrB2-SiC, HfB2-SiC and HfB2-SiC-Y2O3 cylindrical samples were fabricated by SPS and oxidized for 20 s under water vapour atmosphere at 1200, 1550, 1900, 2200 °C and 2400 °C. Up to 1550 °C, only thin oxidation layers are detected. At 1900 °C, each sample develops a silicarich glassy layer. HS and ZS show mechanical damage, while HSY keeps its structural integrity. At 2200 °C, great mechanical issues are observed for ZS, HS exhibits a non-adherent oxide layer, while HSY presents adherent oxide layers, even though the top HfO2 layer was internally split. The observed differences between HS and HSY are explained by: 1) the volume contraction due to the monoclinic to tetragonal transition of HfO2 which is more preponderant in HS than HSY, because Y2O3 stabilizes cubic structure, 2) the creation of a Y2Si2O7 interphase between the HfO2 and the SiC-depleted HfB2 layer that mitigates the thermal mismatch, thus favoring adherence between layers. Overall, ZS has inferior oxidation resistance at high temperatures under water vapour environment, as it cannot develop a stable protective layer that remains adherent to the material beneath, and would likely not sustain high velocity environments. HS and HSY are more promising, but decohesion is observed for HS at 2200 °C and more. At 2400 °C, the top HfO2 layer developed by HS show evidence of sintering, and might provide further resistance towards oxidation, but still exhibits great mechanical damage.  Acknowledgements  The authors would like to thank Claire Sanchez (ONERA) for X-ray diffraction, Dr. P. Beauchêne (ONERA) for the BLOX assistance and fruitful discussions. This work was funded by ONERA.  6 References  [1]  [2]  J. Han, P. Hu, X. Zhang, S. Meng, Oxidation behavior of zirconium diboride-silicon carbide at 1800 °C, Scr. Mater. 57 (2007) 825-828, http://dx.doi.org/10.1016/j. scriptamat.2007.07.009. C.M. Carney, P. Mogilvesky, T.A. Parthasarathy, Oxidation Behavior of Zirconium  431  \\x0c', 'V. Guérineau, A. Julian-Jankowiak  Journal of the European Ceramic Society 38 (2018) 421-432  [25]  [23]  [21]  [20]  [19]  doi.org/10.1111/j.1151-2916.1999.tb02005.x . [18] N. Jacobson, S. Farmer, A. Moore, H. Sayir, High-Temperature oxidation of boron nitride: I, monolithic boron nitride, J. Am. Ceram. 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  "_id": 166,
  "PDF": "Oxidation of Hafnium Carbide and Hafnium Carbide with Additions of Tantalum and Praseodymium.pdf",
  "Text": "['O x i d a t i o n   o f   M e t a l s ,   V o l .   3 6 ,   N o s .   5 / 6 ,   1 9 9 1   O x i d a t i o n   o f   H a f n i u m   C a r b i d e   a n d   H a f n i u m   C a r b i d e   w i t h   A d d i t i o n s   o f   T a n t a l u m   a n d   P r a s e o d y m i u m   E .   L .   C o u r t r i g h t , *   J .   T .   P r a t e r , * *   G .   R .   H o l c o m b , t   G .   R .   S t .   P i e r r e , $   a n d   R .   A .   R a p p $   R e c e i v e d   M a r c h   1 2 ,   1 9 9 1   T h e   o x i d a t i o n   b e h a v i o r   o f   H f C ,   H f C 2 5   w t .   %   T a C ,   a n d   H f C 7   w t .   %   P r C 2   h a s   b e e n   s t u d i e d   b e t w e e n   1 2 0 0 2 2 0 0 ~   P a r a b o l i c   g r o w t h   o f   t h e   o x i d e   l a y e r   h a s   b e e n   o b s e r v e d   f o r   b o t h   H f C   a n d   H f C T a C   o v e r   t h e   e n t i r e   t e m p e r a t u r e   r a n g e .   A   b r e a k   i n   t h e   t e m p e r a t u r e   d e p e n d e n c e   o f   t h e   o x i d a t i o n   k i n e t i c s   o c c u r s   a r o u n d   1 6 0 0 ~   A t   l o w e r   t e m p e r a t u r e s ,   t h e   k i n e t i c s   a r e   l i m i t e d   b y   g a s e o u s   d i f f u s i o n   v i a   p o r e s   i n   t h e   o x i d e .   A b o v e   1 8 0 0 ~   g a s e o u s   d i f f u s i o n   t h r o u g h   p o r e s   b e c o m e s   l e s s   i m p o r t a n t   a s   s c a l e g r o w t h   k i n e t i c s   a r e   d o m i n a t e d   b y   b u l k   ( a m b i p o l a r )   d i f f u s i o n   o f   o x y g e n   a n d   e l e c t r o n s   t h r o u g h   t h e   o x i d e .   K E Y   W O R D S :   h a f n i u m   c a r b i d e ;   c a r b o n   c a r b o n   c o m p o s i t e s ;   o x i d a t i o n   r e s i s t a n c e ;   h i g h   t e m p e r a t u r e .   I N T R O D U C T I O N   C a r b o n c a r b o n   c o m p o s i t e   t e c h n o l o g y   o f f e r s   e x c i t i n g   p o t e n t i a l   f o r   a d v a n c e d   a i r c r a f t t u r b i n e e n g i n e   p e r f o r m a n c e .   T h e   c o m b i n a t i o n   o f   l o w   d e n s i t y   a n d   h i g h   s t r e n g t h   a t   t e m p e r a t u r e s   a b o v e   1 5 0 0 ~   p r o v i d e s   t h e   f l e x i b i l i t y   t o   c o n   s i d e r   n e w   d e s i g n   c o n c e p t s .   H o w e v e r ,   p r o t e c t i n g   c a r b o n   c a r b o n   c o m p o n e n t s   f r o m   t h e   o x i d i z i n g   e n v i r o n m e n t   o f   a   g a s   t u r b i n e   r e p r e s e n t s   t h e   m a j o r   c h a l   l e n g e   t o   r e a l i z i n g   t h i s   p o t e n t i a l .   * P a c i f i c   N o r t h w e s t   L a b o r a t o r y ,   R i c h l a n d ,   W a s h i n g t o n .   * * P r e s e n t l y   a t   A r m y   R e s e a r c h   O f f i c e ,   R e s e a r c h   T r i a n g l e   P a r k ,   N o r t h   C a r o l i n a .   t P r e s e n t l y   a t   U . S .   B u r e a u   o f   M i n e s ,   A l b a n y ,   O r e g o n .   $ O h i o   S t a t e   U n i v e r s i t y ,   C o l u m b u s ,   O h i o .   4 2 3   0 0 3 0 7 7 0 X / 9 1 / 1 2 0 0 0 4 2 3 5 0 6 . 5 0 / 0   (cid:14) 9   1 9 9 1   P l e n u m   P u b l i s h i n g   C o r p o r a t i o n   \\x0c', \"4 2 4   C o u r t r i g h t   e t   a l .   C u r r e n t l y ,   c a r b o n c a r b o n   c o m p o s i t e s   r e l y   o n   s i l i c o n c a r b i d e   s u r f a c e   l a y e r s ,   d e p o s i t e d   e i t h e r   b y   c h e m i c a l   v a p o r   d e p o s i t i o n   o r   p a c k   c e m e n t a t i o n ,   t o   p r o v i d e   o x i d a t i o n   r e s i s t a n c e .   T h i s   c o m b i n a t i o n   w o r k s   w e l l   f o r   t e m p e r a   t u r e s   u p   t o   a p p r o x i m a t e l y   1 7 6 0 ~   F o r   p r o t e c t i o n   a b o v e   1 7 6 0 ~   a   n e w   p r o t e c t i v e   s y s t e m   i s   n e e d e d .   F o r   t h e r m o d y n a m i c   r e a s o n s ,   t h e   f i r s t   l a y e r   t o   b e   p l a c e d   i n   c o n t a c t   w i t h   t h e   c a r b o n c a r b o n   c o m p o s i t e   w i l l   p r o b a b l y   h a v e   t o   b e   a   c a r b i d e   t h a t   u p o n   o x i d a t i o n   f o r m s   a n   o x i d e   l a y e r   w h i c h   c a n   p r o v i d e   a t   l e a s t   m i n i m a l   o x i d a t i o n   r e s i s t a n c e .   R e l a t i v e l y   f e w   r e f r a c t o r y   o x i d e s   a r e   s t a b l e   i n   a n   o x i d i z i n g   a t m o s p h e r e   a b o v e   1 7 6 0 ~   O n l y   t h o r i u m ,   h a f n i u m ,   z i r c o n i u m ,   a n d   b e r y l l i u m   h a v e   s u f f i c i e n t l y   h i g h   m e l t i n g   p o i n t s   t o   b e   c o n s i d e r e d   a s   c a n d i d a t e   b a s e   c o n s t i t   u e n t s   f o r   t h e   o x i d a t i o n r e s i s t a n t   c a r b i d e   l a y e r .   O f   t h e s e ,   t h o r i u m   i s   r a d i o a c   t i v e   a n d   b e r y l l i u m   i s   p o t e n t i a l l y   t o x i c ;   t h u s   z i r c o n i u m   a n d   h a f n i u m   c a r b i d e   ( H f C )   a r e   p r e f e r a b l e   c h o i c e s   f o r   a   h i g h t e m p e r a t u r e   c o a t i n g   s y s t e m .   H a f   n i u m   o x i d e   h a s   b e t t e r   h i g h t e m p e r a t u r e   s t a b i l i t y   a n d   m a y   b e   t h e   b e t t e r   d i f f u s i o n   b a r r i e r   1   d e s p i t e   i t s   r e l a t i v e l y   h i g h   o x y g e n   m o b i l i t y .   T h e   o x i d a t i o n   o f   H f C   h a s   b e e n   s t u d i e d   t o   1 2 0 0 ~   2 ' 3   b u t   t h e   o x i d e g r o w t h   k i n e t i c s   h a v e   n o t   b e e n   u n a m b i g u o u s l y   e s t a b l i s h e d .   I n   t h i s   w o r k ,   t h e   o x i d a t i o n   b e h a v i o r   o f   H f C   u p   t o   2 2 0 0 ~   w a s   e x a m i n e d   a n d   t h e   f e a s i b i l i t y   o f   a l l o y i n g   H f C   w i t h   2 5   w t . %   t a n t a l u m   c a r b i d e   ( T a C )   a n d   1 0   w t . %   p r a s e o d y m i u m   c a r b i d e   ( P r C 2 )   t o   r e d u c e   o x y g e n   p e r m e a b i l i t y   i n   t h e   g r o w i n g   o x i d e   l a y e r   w a s   i n v e s t i g a t e d .   O x i d a t i o n   s t u d i e s   o n   t h e   h a f n i u m t a n t a l u m   m e t a l l i c   a l l o y   s y s t e m   h a v e   s h o w n   t h a t   h a f n i u m   a l l o y s   w i t h   2 0 3 0   w t . %   t a n t a l u m   a r e   s i g n i f i c a n t l y   m o r e   o x i d a t i o n   r e s i s t a n t   t h a n   p u r e   h a f n i u m .   I n   i s o t h e r m a l   t e s t i n g   a t   2 2 0 0 ~   t h e   r a t e   o f   o x i d e   g r o w t h   w a s   h a l f   t h a t   o f   p u r e   h a f n i u m ,   4   a n d   t a n t a l u m   p r o m o t e d   t h e   f o r m a t i o n   o f   a   d e n s e ,   t e n a c i o u s   o x i d e   l a y e r .   C e r t a i n   r a r e e a r t h   e l e m e n t s ,   f o r   e x a m p l e   p r a s e o d y m i u m ,   c a n   b e   a d d e d   t o   h a f n i a   t o   s t a b i l i z e   t h e   p y r o c h l o r e   s t r u c t u r e   H f 2 P r 2 0 7 .   T h e   p y r o c h l o r e s   h a v e   a n i o n i c   m o b i l i t i e s   t h a t   a r e   a n   o r d e r   o f   m a g n i t u d e   l o w e r   t h a n   t h a t   o f   t h e   h i g h t e m p e r a t u r e   f l u o r i t e   p h a s e   5   a n d   h a v e   m e l t i n g   t e m p e r a t u r e s   a b o v e   2 3 0 0 ~   T h u s ,   i f   a n   o x i d e   o f   t h e   p y r o c h l o r e   c a n   b e   g r o w n   f r o m   a   m i x e d   c a r b i d e ,   i t   m i g h t   p r o v i d e   b e t t e r   o x i d a t i o n   p r o t e c t i o n   t h a n   a   p u r e   H f O 2   l a y e r .   E X P E R I M E N T A L   P R O C E D U R E S   O x i d a t i o n   s t u d i e s   w e r e   p e r f o r m e d   o n   c a r b i d e   s p e c i m e n s   c u t   f r o m   l a r g e r   h o t p r e s s e d   c o m p a c t s   ( 3 . 1   c m   d i a m e t e r   x   2 . 6   c m ) .   T h e   h o t   p r e s s i n g   w a s   d o n e   a t   L o s   A l a m o s   N a t i o n a l   L a b o r a t o r y   u s i n g   p r e h o m o g e n i z e d   p o w d e r s   p r o c u r e d   f r o m   t h e   C e r a c   C o r p o r a t i o n .   T h e   p r e s s i n g   c o n d i t i o n s   a n d   p h y s i c a l   p a r a m e t e r s   f o r   t h e   t h r e e   c a r b i d e   c o m p a c t s   a r e   l i s t e d   i n   T a b l e   I .   X r a y   d i f f r a c t i o n   p a t t e r n s   o f   a l l   t h r e e   c o m p a c t s   i n d i c a t e d   t h a t   t h e   m i x e d   \\x0c\", 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   T a b l e   I .   C a r b i d e   S p e c i m e n s   4 2 5   C o m p o s i t i o n   a   P r e s s i n g   X r a y   ( w t . % )   c o n d i t i o n s   D e n s i t y   l a t t i c e   c o n s t a n t   H f C   2   (cid:141)   1 0 7   P a / 3 0 0 0 ~   1 1 . 6   g m / c m   3   4 . 6 3 8   A   ( 9 1 %   t h e o r e t i c a l )   ( s t o i c h i o m e t r i c   c a r b i d e )   H f C   2 5   T a C   2   x   1 0 7   P a / 2 7 5 0 ~   1 1 . 9   g m / c m   3   4 . 6 0 6   A   ( 9 0 %   t h e o r e t i c a l )   ( s t o i c h i o m e t r i c   c a r b i d e )   H f C   7   P r C 2   2   x   1 0   v   P a / 2 2 0 0 ~   1 0 . 3   g m / c m   3   4 . 6 3 0   A   ( 8 4 %   t h e o r e t i c a l )   a T y p i c a l   i m p u r i t i e s :   0 . 1 %   A t ,   T i ,   Z n ,   a n d   Z r ,   0 . 0 5 %   F e ,   C u ,   N i ,   a n d   S i .   5   U 3   { 3   T e s t   C h a m b e r   F   (cid:1) 8 9   R u p l : u r e   D i s k   E x h a u s t   W i n d o ~   F i g .   1 .   L a s e r h e a t i n g   a p p a r a t u s   u s e d   f o r   h i g h t e m p e r a t u r e   o x i d a t i o n   o f   c a r b i d e s .   \\x0c', '4 2 6   C o u r t r i g h t   e t   a L   c a r b i d e s   h a d   f o r m e d   h o m o g e n e o u s   s o l i d   s o l u t i o n s .   T h e   H f C   a n d   H f C   T a C   c o m p a c t s   w e r e   p r e s s e d   t o   f i n a l   d e n s i t i e s   n e a r   9 0 %   ( t h e o r e t i c a l ) .   T h e   p r e s s i n g   o f   t h e   H f C P r C 2   p o w d e r   w a s   s u s p e n d e d   p r e m a t u r e l y   a t   2 2 0 0 ~   p r i o r   t o   s i g n i f i c a n t   h e a d   m o v e m e n t   i n   t h e   d i e .   A   p r e s s u r e   s p i k e   a t t r i b u t e d   t o   t h e   d e c o m p o s i t i o n   o f   t h e   P r C 2   a n d   t h e   r e l e a s e   o f   p r a s e o d y m i u m   v a p o r   w a s   o b s e r v e d   i n   t h e   s y s t e m .   T h e   h o t p r e s s i n g   o p e r a t i o n   w a s   t h e r e f o r e   h a l t e d   t o   p r e v e n t   f u r t h e r   l o s s   o f   p r a s e o d y m i u m .   T h e   r e s u l t   w a s   a   s a m p l e   w i t h   a   f i n a l   d e n s i t y   o f   a b o u t   8 4 %   ( t h e o r e t i c a l )   a n d   a   f i n a l   c o m p o s i t i o n   o f   7   w t . %   p r a s e o d y m i u m .   O x i d a t i o n   s t u d i e s   w e r e   p e r f o r m e d   a t   1 2 0 0 1 5 3 0 ~   u s i n g   t h e r m o g r a v i   m e t r i c   a n a l y s i s   ( T G A )   a n d   a t   1 8 0 0 2 2 0 0 ~   u s i n g   a   l a s e r h e a t i n g   a p p a r   a t u s .   T h e   T G A   e x p e r i m e n t s   w e r e   c o n d u c t e d   i n   a i r   b y   p r e h e a t i n g   a   f u r n a c e   t o   t h e   d e s i r e d   t e m p e r a t u r e   a n d   t h e n   l o w e r i n g   t h e   r e c t a n g u l a r   s a m p l e   ( 1 5   x   8   x   1   m m )   i n t o   t h e   h o t   z o n e .   W e i g h t c h a n g e s   f o r   t h e   s p e c i m e n s   w e r e   r e c o r d e d   a s   a   f u n c t i o n   o f   t i m e ,   a n d   t h e   w e i g h t g a i n   d a t a   w e r e   t h e n   c o n v e r t e d   t o   o x i d e   t h i c k n e s s   b y   a s s u m i n g   a   1 0 0 %   d e n s e   o x i d e   l a y e r .   A d d i t i o n a l   e x p e r i m e n t s   o n   H f C   w e r e   d o n e   w i t h   o x y g e n   p a r t i a l   p r e s s u r e s   o f   0 . 0 2   a n d   1   a t .   O x i d a t i o n   o f   t h e   c a r b i d e s   a b o v e   1 8 0 0 ~   w a s   p e r f o r m e d   b y   u s i n g   a   d e f o c u s e d   C O 2   l a s e r   t o   h e a t   t h e   s a m p l e .   O x i d a t i o n   k i n e t i c s   w e r e   e s t a b l i s h e d   b y   e x p o s i n g   s a m p l e s   t o   f l o w i n g   a i r   f o r   v a r y i n g   t i m e s   a t   a   g i v e n   t e m p e r a t u r e   a n d   t h e n   m e t a l l o g r a p h i c a l l y   e x a m i n i n g   e a c h   s a m p l e   t o   d e t e r m i n e   t h e   t h i c k   n e s s   o f   t h e   o x i d e   l a y e r s   t h a t   f o r m e d   o n   t h e   l a s e r h e a t e d   s i d e   o f   t h e   s a m p l e .   T h e   t e m p e r a t u r e   o n   t h e   f r o n t   s u r f a c e   o f   t h e   s a m p l e   w a s   m o n i t o r e d   w i t h   a   t w o c o l o r   i n f r a r e d   p y r o m e t e r ,   a n d   t h e   o u t p u t   w a s   f e d   b a c k   t o   a   c o m p u t e r   w h i c h   c o n t r o l l e d   t h e   p o w e r   t o   t h e   l a s e r .   I n   t h i s   m a n n e r ,   t h e   t e m p e r a t u r e   d u r i n g   t h e   o x i d a t i o n   e x p o s u r e   c o u l d   b e   c o n t r o l l e d   t o   f o l l o w   a   p r e d e t e r m i n e d   t h e r m a l   h i s t o r y .   A   s c h e m a t i c   o f   t h e   o x i d a t i o n   a p p a r a t u s   i s   p r e s e n t e d   i n   F i g .   1 .   E x p e r i m e n t s   w e r e   p e r f o r m e d   o n   d i s c   s p e c i m e n s   ( 3 . 8 m m   d i a .   x   0 . 9   m m   t h i c k )   t h a t   h a d   b e e n   m a c h i n e d   f r o m   t h e   h o t p r e s s e d   c a r b i d e   s a m p l e s .   R E S U L T S   P a r a b o l i c   g r o w t h   k i n e t i c s   w e r e   o b s e r v e d   f o r   H f C   a n d   H f C T a C   o v e r   t h e   e n t i r e   t e m p e r a t u r e   r a n g e .   T h e s e   g r o w t h   k i n e t i c s   w e r e   d e s c r i b e d   b y   t h e   e x p r e s s i o n :   R = k p t l / 2   +   R o   ( 1 )   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   4 2 7   w h e r e   R   i s   a   k i n e t i c   p a r a m e t e r   i n d i c a t i n g   t h e   e x t e n t   o f   r e a c t i o n   ( s a m p l e   w e i g h t g a i n   o r   o x i d e   t h i c k n e s s ) ,   k p   t h e   p a r a b o l i c   r a t e   c o n s t a n t ,   t   i s   i s o t h e r m a l   r e a c t i o n   t i m e ,   a n d   R o   t h e   k i n e t i c   p a r a m e t e r   a t   t   =   0 .   H f C   T h e   o x i d e   l a y e r   t h a t   f o r m e d   o n   H f C   d i s p l a y e d   p a r a b o l i c   g r o w t h   k i n e t i c s   f o r   b o t h   t e m p e r a t u r e   r e g i m e s .   T h e   t e m p e r a t u r e   d e p e n d e n c e   o f   t h e   o x i d e   g r o w t h   c o n s t a n t   f o r   H f C   i s   p r e s e n t e d   i n   F i g .   2 .   T h e r e   i s   a   b r e a k   i n   t h e   c u r v e   a t   a b o u t   1 8 0 0 ~   B e l o w   1 6 0 0 ~   t h e   g r o w t h   c o n s t a n t   ( k p )   w a s   p r o p o r t i o n a l   t o   0 . 0 0 3 2   p ~   a n d   c o u l d   b e   d e s c r i b e d   b y   t h e   e q u a t i o n :   k p = 4 . 8 3   x   1 0   4   e x p ( 7 1 6 / T ( K ) )   c m / s   1 / 2   ( T =   1 2 0 0 1 8 0 0 ~   ( 2 )   ( T =   1 4 7 3 2 0 7 3   K )   \" 7   c   O   . _ o   g   O   I 1 )   n   2   1 0 0   4 ( 1     2 O   1 0     8     4     2     1   2 . 5   ~ m p e r a t u r e   ( C )   o   o   o o ~   o   o   o   o   o o o   o   o   o   I   I \\\\ [ 1 1 1   I   \\\\ ~ ,   ( R e f .   8 )   %   \\\\   \\\\   \\\\   I   I   I   I   I   I   I   ~ \\' ,   I   3 . 0   3 . 5   4 . 0   4 . 5   5 . 0   5 . 5   6 . 0   6 . 5   1 0 , 0 0 0 / T e m p e r a t u r e   ( K )   7 . 0   F i g .   2 .   P a r a b o l i c   r a t e   c o n s t a n t   f o r   o x i d e   g r o w t h   o n   H f C ,   a n d   f o r   c o m p a r i s o n ,   d a t a   f o r   p u r e   H f   ( R e f .   8 )   e x t r a p o l a t e d   f r o m   9 0 0 1 2 0 0 ~   \\x0c', 'O 0   F i g .   3 .   M i c r o g r a p h   o f   H f C   s a m p l e   o x i d i z e d   f o r   5   m i n   a t   2 0 0 0 ~   r   o   e ~   \\x0c', \"O x i d a t i o n   o f   H a f n i u m   C a r b i d e   4 2 9   A b o v e   1 8 0 0 ~   t h e   t e m p e r a t u r e   d e p e n d e n c e   o f   t h e   i s o t h e r m a l   g r o w t h   c o n   s t a n t   i s   d e s c r i b e d   b y :   k p   =   0 . 5 0   e x p (   1 3 7 0 0 / T ( K ) )   c m / s   1 / 2   ( T =   1 8 0 0 2 2 0 0 ~   ( 3 )   ( T =   2 0 7 3 2 4 7 3   K )   T h e   o x i d e   l a y e r   t h a t   f o r m e d   w a s   p e r m e a t e d   w i t h   s m a l l   p o r e s   a n d   c r a c k s .   T h e   d i s t r i b u t i o n   o f   t h e s e   p o r e s   a n d   c r a c k s   v a r i e d   a c r o s s   t h e   t h i c k n e s s   o f   t h e   o x i d e   l a y e r   ( s e e   F i g .   3 ) .   A t   t h e   c a r b i d e o x i d e   i n t e r f a c e   t h e r e   w a s   a   n a r r o w   b a n d   ( a b o u t   1 0   ~ m   w i d e )   o f   o x i d e   t h a t   w a s   r e l a t i v e l y   f e a t u r e l e s s ,   w i t h   t h e   e x c e p t i o n   o f   a   f e w   i n t e r m e d i a t e   s i z e   p o r e s   ( r e m n a n t s   o f   v o i d s   p r e s e n t   i n   t h e   c o m p a c t ) .   I n   c o n t r a s t ,   t h e   m i c r o s t r u c t u r e   o f   t h e   o u t e r   o x i d e   l a y e r   w a s   c h a r a c t e r i z e d   b y   a   d e n s e   p o p u l a t i o n   o f   s u b m i c r o n   p o r e s   a n d   a   m o d e r a t e   a m o u n t   o f   o x i d e   c r a c k i n g .   I n   t h e   p r e s e n t   s t u d y ,   t h e   o x i d e   l a y e r   t h a t   f o r m e d   o n   t h e   H f C   w a s   H f O 2 .   X r a y   d i f f r a c t i o n   p a t t e r n s   i n d i c a t e d   t h a t   t h e   o x i d e   w a s   m o n o c l i n i c   a t   a m b i e n t   t e m p e r a t u / e .   H o w e v e r ,   r e f e r e n c e   t o   t h e   p h a s e   d i a g r a m   b y   R u h   a n d   P a t e l   6   s u g g e s t s   t h a t   t h e   o x i d e   t h a t   f o r m e d   o n   s a m p l e s   e x p o s e d   a b o v e   a b o u t   1 7 0 0 ~   m u s t   h a v e   g r o w n   a s   t e t r a g o n a l   H f O 2   a n d   t r a n s f o r m e d   t o   t h e   m o n o c l i n i c   p h a s e   u p o n   c o o l i n g .   T h e r e   i s   a   m o d e s t   v o l u m e   e x p a n s i o n ,   1 2 . 7 % ,   7   a s s o c i a t e d   w i t h   t h i s   t r a n s f o r m a t i o n .   H o w e v e r ,   c o m p a r i s o n   o f   t h e   c r a c k i n g   p a t t e r n s   o n   s a m p l e s   o x i d i z e d   a b o v e   a n d   b e l o w   t h e   t r a n s f o r m a t i o n   t e m p e r a t u r e   d o   n o t   a p p e a r   t o   d i f f e r   s i g n i f i c a n t l y ,   s u g g e s t i n g   t h a t   t h e   t e t r a g o n a l t o m o n o c l i n i c   t r a n s f o r m a t i o n   i s   n o t   a   m a j o r   r e a s o n   f o r   t h e   o x i d e   c r a c k i n g .   A t   1 7 0 0 ~   t h e   o x i d e   i s   q u i t e   p l a s t i c   a n d   i s   a p p a r e n t l y   a b l e   t o   a c c o m m o d a t e   t h e   t r a n s f o r m a   t i o n   s t r e s s e s   t h r o u g h   d e f o r m a t i o n ,   b u t   i t   a p p e a r s   t h a t   t h e   m a j o r i t y   o f   t h e   c r a c k i n g   w a s   p r o d u c e d   b y   t h e r m a l   ' s t r e s s e s   g e n e r a t e d   a s   t h e   s a m p l e s   w e r e   c o o l e d   d o w n   b e l o w   1 0 0 0 ~   H f C T a C   T h e   o x i d e   l a y e r s   t h a t   f o r m e d   o n   H f C T a C   a l s o   o b e y e d   p a r a b o l i c   g r o w t h   k i n e t i c s .   T h e   t e m p e r a t u r e   d e p e n d e n c e   o f   t h e   g r o w t h   c o n s t a n t   i s   p r e   s e n t e d   i n   F i g .   4 .   A g a i n ,   a   b r e a k   i n   t h e   k i n e t i c s   o c c u r s   a r o u n d   1 8 0 0 ~   B e l o w   1 6 0 0 ~   t h e   i s o t h e r m a l   g r o w t h   c o n s t a n t   i s   d e s c r i b e d   b y   t h e   e q u a t i o n :   k p   =   6 . 2 9   x   1 0   4   e x p ( 1 3 9 0 / T ( K ) )   c m / s   1 / 2   ( T   =   1 2 0 0 1 8 0 0 ~   ( 4 )   A b o v e   1 8 0 0 ~   t h e   g r o w t h   c o n s t a n t   i s   g i v e n   b y :   k p   =   0 . 1 3   e x p ( 9 5 0 0 / T ( K ) )   c m / s   1 / 2   ( T >   1 8 0 0 ~   ( 5 )   T h e   s u r f a c e   o f   t h e   o x i d e   t h a t   f o r m e d   a t   2 2 0 0 ~   w a s   v e r y   i r r e g u l a r ,   m a k i n g   r e l i a b l e   t h i c k n e s s   m e a s u r e m e n t s   i m p o s s i b l e .   I t   i s   b e l i e v e d   t h a t   r a p i d   C O   g a s   \\x0c\", '4 3 0   C o u r t r i g h t   e t   a l .   0   , _ o   \" O   \" 5   I D   r r   c   o   ( . 9   1 0 0   8 0   4 0   2 0     1 0     8     4     2     1   2 . 5   ~ m p e r a t u r e   ( C )   _ 1   I I I I I   I   \" , ~ ,   H f 6   T a   a l l o y   \\\\   , , ,   o   \\\\   H f C + T a C   H f 2 5   T a   a l l o y   I   I   I   I   I   I   I   I   3 . 0   3 . 5   4 , 0   4 . 5   5 , 0   5 . 5   6 . 0   6 . 5   1 0 , 0 0 0 / T e m p e r a t u r e   ( K )   7 . 0   F i g .   4 .   P a r a b o l i c   r a t e   c o n s t a n t   f o r   o x i d e   g r o w t h   o n   H f C 2 5   T a C .   e v o l u t i o n   m a y   h a v e   d i s r u p t e d   t h i s   l a y e r   d u r i n g   t h e   e a r l y   s t a g e s   o f   t h e   o x i d e   g r o w t h .   T h e   H f C T a C   o x i d a t i o n   k i n e t i c s   a b o v e   1 8 0 0 ~   w e r e   c o m p a r a b l e   t o   t h o s e   r e p o r t e d   p r e v i o u s l y   f o r   t h e   h a f n i u m t a n t a l u m   m e t a l l i c   a l l o y ,   4   a l s o   s h o w n   i n   F i g .   4 .   A   d e n s e ,   t e n a c i o u s ,   c r a c k r e s i s t a n t   o x i d e   f o r m s   o n   t h e   h a f n i u m   t a n t a l u m   a l l o y   w h e r e a s   t h e   o x i d e   t h a t   f o r m s   o n   t h e   c a r b i d e   i s   p o r o u s   a n d   p r o n e   t o   c r a c k i n g .   T h e   p o r o u s   o x i d e   s c a l e   f o r m e d   o n   t h e   c a r b i d e s   r e s u l t s   f r o m   p o r e   g e n e r a t i o n   i n d u c e d   b y   e s c a p i n g   C O ,   a   b y p r o d u c t   o f   c a r b i d e   o x i d a t i o n .   F o r   t h e   m e t a l l i c   a l l o y s ,   i n t e r n a l   o x i d a t i o n   o f   t h e   h a f n i u m   i n   t h e   b a s e   m e t a l   r e s u l t s   i n   a   l a m e l l a r   s t r u c t u r e .   T h e   H f O 2   p l a t e l e t s   e x t e n d   i n t o   t h e   u n d e r l y i n g   m e t a l ,   h e l p   t o   a n c h o r   t h e   o x i d e   t o   t h e   b a s e   a l l o y ,   a n d   s e r v e   a s   a   m e c h a n i c a l   s u p e r s t r u c t u r e   t o   t h e   s c a l e .   N o   s i m i l a r   m e c h a n i s m   o c c u r s   f o r   t h e   m i x e d   c a r b i d e .   I n   c o m p a r i s o n ,   t h e   o x i d e   t h a t   f o r m e d   o n   t h e   c a r b i d e   i s   s i g n i f i c a n t l y   w e a k e n e d   b y   t h e   p r e s e n c e   o f   p o r e s   w h i c h   n e c e s s a r i l y   p e r m i t   C O   e v o l u t i o n .   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   4 3 1   B e l o w   1 8 0 0 ~   t h e   o x i d e   l a y e r   t h a t   f o r m s   o n   t h e   H f C T a C   s a m p l e s   w a s   H f O 2 ;   b u t   a b o v e   1 8 0 0 ~   t h e   o x i d a t i o n   p r o d u c t   i s   a   m i x t u r e   o f   H f O 2   a n d   T a 2 H f 6 0 1 9 .   T h e   m ~ c r o s t r u c t u r e s   o f   t h e   o x i d e s   o n   t h e   H f C   T a C   s a m p l e s   w e r e   s i m i l a r   t o   t h o s e   o f   t h e   H f C ,   e x c e p t   t h a t   t h e   o x i d e   w a s   m o r e   h e a v i l y   c r a c k e d   a n d   t h e   c r a c k s   e x t e n d e d   t h r o u g h o u t   t h e   e n t i r e   t h i c k n e s s   o f   t h e   o x i d e .   T h e   u n d e r l y i n g   c a r b i d e   w a s   h e a v i l y   c r a c k e d   t o   a   d e p t h   o f   a b o u t   2 0 / 3 m .   I n   s a m   p l e s   o x i d i z e d   a t   2 0 0 0 ~   t h e   o u t e r   o x i d e   s u r f a c e   w a s   v e r y   r a g g e d ,   s u g g e s t i n g   t h a t   t h e   e v o l u t i o n   o f   C O   g a s   b u b b l e s   m a y   h a v e   d i s r u p t e d   t h e   g r o w t h   o f   t h e   o x i d e   l a y e r   d u r i n g   t h e   e a r l y   s t a g e s   o f   o x i d a t i o n .   H f C P r C 2   N o   r a t e   k i n e t i c s   w e r e   o b t a i n e d   f o r   H f C P r C 2   b e c a u s e   t h e   m a t e r i a l   d i s   p l a y e d   r u n a w a y   o x i d a t i o n .   D u r i n g   o x i d a t i o n ,   t h e   H f C P r C 2   g e n e r a t e d   e x c e s   s i v e   h e a t   w h i c h   m a d e   t e m p e r a t u r e   c o n t r o l   i m p o s s i b l e .   A b o v e   1 2 0 0 ~   t h e   o x i d e   l a y e r   p e r i o d i c a l l y   d i s i n t e g r a t e d   i n t o   a   f i n e   d u s t .   X r a y   d i f f r a c t i o n   p a t t e r n s   o f   t h e   o x i d e   i n d i c a t e d   t h e   p r e s e n c e   o f   b o t h   m o n o c l i n i c   a n d   c u b i c   ( e i t h e r   t h e   f l u o r i t e   o r   p y r o c h l o r e )   H f O 2   i n   r o u g h l y   e q u a l   a m o u n t s .   T h e   f l u o r i t e   a n d   p y r o c h l o r e   s t r u c t u r e s   a r e   v e r y   s i m i l a r   a n d   p r o d u c e   s i m i l a r   d i f f r a c t i o n   p a t t e r n s ,   b u t   t h e   p y r o c h l o r e   p a t t e r n   h a s   a   f e w   a d d i t i o n a l   l o w   i n t e n s i t y   p e a k s .   T h e   d i f f r a c t i o n   p a t t e r n s   f r o m   o u r   s a m p l e s   w e r e   t o o   w e a k   t o   d i s t i n g u i s h   c o n c l u s i v e l y   b e t w e e n   t h e   t w o   s t r u c t u r e s ,   a l t h o u g h   t h e   p r a s e o d y m i u m   c o n c e n t r a t i o n s   i n   t h e   o x i d e   v a r i e d   f r o m   a b o u t   1 0 %   f o r   s a m p l e s   o x i d i z e d   a t   1 8 0 0 ~   a n d   b e l o w ,   t o   l e s s   t h a n   5 %   f o r   s a m p l e s   o x i d i z e d   a t   2 2 0 0 ~   B a s e d   o n   t h e s e   r e l a t i v e l y   l o w   p r a s e o d y m i u m   c o n c e n t r a   t i o n s ,   t h e   c u b i c   H f O 2   h a s   t e n t a t i v e l y   b e e n   i d e n t i f i e d   a s   t h e   f l u o r i t e   p h a s e ,   a n d   n o t   t h e   p y r o c h l o r e ,   w h i c h   w o u l d   r e q u i r e   m u c h   h i g h e r   c o n c e n t r a t i o n s   o f   p r a s e o d y m i u m   t o   b e   s t a b l e .   I n   s e v e r a l   e x p e r i m e n t s ,   t h e   o x i d a t i o n   w a s   i n t e r r u p t e d   p r i o r   t o   c o m p l e t e   c o n s u m p t i o n   o f   t h e   c a r b i d e   b y   e x c l u d i n g   a i r   f r o m   t h e   s a m p l e .   M e t a l l o   g r a p h i c   c r o s s s e c t i o n s   o f   t h e s e   s a m p l e s   w e r e   p r e p a r e d   a n d   e x a m i n e d   w i t h   a   s c a n n i n g   e l e c t r o n   m i c r o s c o p e   e q u i p p e d   w i t h   e n e r g y   d i s p e r s i v e   x r a y   a n a l y s i s .   T h i s   r e v e a l e d   t h e   p r e s e n c e   o f   a n   i n t e r n a l l y   o x i d i z e d   p h a s e   o f   p u r e   P r 2 0 3   c o n c e n t r a t e d   a l o n g   t h e   i n t e r n a l   p o r o s i t y   o f   t h e   c a r b i d e   s a m p l e ,   s i m i l a r   t o   o b s e r v a t i o n s   i n   a l l o y   s y s t e m s   w h e r e   t h e   c o n c e n t r a t i o n   o f   a n   a c t i v e   s p e c i e s   i s   t o o   l o w   t o   p e r m i t   t h e   g r o w t h   o f   a   c o n t i n u o u s   s u r f a c e   o x i d e   l a y e r .   T h u s ,   i n t e r n a l   o x i d a t i o n   o f   p r a s e o d y m i u m ,   a n d   v o l a t i l i z a t i o n   o f   t h e   m e t a l   a t   t h e   h i g h e r   t e m p e r a t u r e s ,   p r e v e n t e d   t h e   p r a s e o d y m i u m   f r o m   d i f f u s i n g   i n t o   t h e   g r o w i n g   o x i d e   l a y e r .   T h e r m a l   C y c l i n g   T h e   H f C   a n d   H f C T a C   s a m p l e s ,   s h o w n   i n   F i g .   5 ,   w e r e   s u b j e c t e d   t o   a   s e r i e s   o f   f i v e   t h e r m a l   c y c l e s   ( i n   a i r )   b e t w e e n   2 0 0 0 ~   a n d   5 0 0 ~   E a c h   c a r b i d e   \\x0c', '1 , 4   F i g .   5 .   M i c r o g r a p h s   o f   H f C   ( l e f t )   a n d   H f C   T a C   ( r i g h t )   f o l l o w i n g   f i v e   t h e r m a l   c y c l e s   b e t w e e n   2 0 0 0 ~   a n d   5 0 0 ~   f l t Q   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   C a r b i d e   4 3 3   s p e c i m e n   w a s   h e l d   a t   2 0 0 0 ~   f o r   a   t o t a l   o f   1 0   m i n   ( 5   x   2   m i n   e x p o s u r e s ) .   H e a t i n g   a n d   c o o l i n g   r a t e s   w e r e   a b o u t   3 0 0 ~   T h e   H f C   s a m p l e   w a s   n o t   n o t i c e a b l y   a f f e c t e d   b y   t h e   t h e r m a l   c y c l i n g   a n d   d e v e l o p e d   a n   o x i d e   l a y e r   t h a t   w a s   s i m i l a r   t o   t h a t   f o r m e d   d u r i n g   a   s i n g l e   i s o t h e r m a l   e x p o s u r e   t o   a i r .   I n   c o n t r a s t ,   t h e   H f C   T a C   s a m p l e   w a s   c l e a r l y   d a m a g e d   b y   t h e   t h e r m a l   c y c l e s .   T h e   o x i d e   w a s   c r a c k e d   a n d   t h e   g r o w t h   k i n e t i c s   w e r e   a p p r o x i m a t e l y   t w i c e   t h o s e   o f   t h e   e q u i v a l e n t   i s o t h e r m a l   e x p o s u r e .   D I S C U S S I O N   I n   c e r t a i n   r e s p e c t s ,   t h e   o x i d a t i o n   b e h a v i o r   o f   t h e   c a r b i d e   s a m p l e s   w a s   s i m i l a r   t o   t h a t   o b s e r v e d   f o r   m e t a l l i c   a l l o y s .   F o r   e x a m p l e ,   p r o t e c t i v e   o x i d e   l a y e r s   d i s p l a y e d   p a r a b o l i c   g r o w t h   k i n e t i c s .   A l s o ,   e v i d e n c e   o f   i n t e r n a l   o x i d a   t i o n   f o r   T a C   a n d   P r C 2   i s   c o m m o n   i n   m e t a l l i c   s y s t e m s   w h e r e   t h e   a d d i t i o n   o f   a   d i l u t e   a c t i v e   a l l o y i n g   e l e m e n t   i s   i n s u f f i c i e n t   t o   p r o m o t e   g r o w t h   o f   a   c o n   t i n u o u s   l a y e r .   D u r i n g   t h e   o x i d a t i o n   o f   a   c a r b i d e ,   g a s e o u s   C O   i s   p r o d u c e d   a s   a   b y   p r o d u c t .   S i n c e   C O   d i s s o l u t i o n   i n   t h e   o x i d e   l a y e r   i s   v i r t u a l l y   n i l ,   t h e   g e n e r a   t i o n   o f   c o n t i n u o u s   p o r o s i t y   i n   t h e   o x i d e   r e p r e s e n t s   t h e   p r i m a r y   m e a n s   f o r   e s c a p e   o f   t h e   g a s .   T h e r e f o r e ,   C O   g e n e r a t i o n   i n t r o d u c e s   p o r o s i t y   w h i c h   a l l o w s   g a s e o u s   d i f f u s i o n   v i a   p o r e s .   G a s   f o r m e d   b e l o w   t h e   o x i d e   l a y e r   c a n   a l s o   l i f t   a n d   d i s r u p t   t h e   o x i d e   l a y e r .   I n   t h e   p r e s e n t   s t u d y ,   t h i s   o c c u r r e d   f o r   t h e   H f C   T a C   a b o v e   2 0 0 0 ~   T h e   g r o w t h   k i n e t i c s   c u r v e s   i n   F i g s .   2   a n d   4   s h o w   a n   u n u s u a l   d i s c o n t i n u   o u s   t e m p e r a t u r e   d e p e n d e n c e   s u g g e s t i n g   t h a t   t h e r e   i s   a   c h a n g e   i n   t h e   r a t e   c o n t r o l l i n g   m e c h a n i s m   a r o u n d   1 8 0 0 ~   S o l i d s t a t e d i f f u s i o n c o n t r o l l e d   k i n e t i c s   a r e   b e l i e v e d   t o   d o m i n a t e   a b o v e   1 8 0 0 ~   s i n c e   t h e   t e m p e r a t u r e   d e p e n d e n c e   f o r   t h e   o x i d e   g r o w t h   r a t e   i s   c l o s e   t o   t h e   o x i d e   g r o w t h   r a t e s   o b s e r v e d   f o r   m e t a l l i c   h a f n i u m   8   a n d   h a f n i u m t a n t a l u m   a l l o y s .   4   T h e   l o w e r   t e m p e r a t u r e   k i n e t i c s   c a n   b e   e x p l a i n e d   b y   a s s u m i n g   t h a t   o x i d e   g r o w t h   i s   c o n t r o l l e d   b y   0 2 ,   C O ,   a n d   C O 2   g a s e o u s   t r a n s p o r t   t h r o u g h   t h e   i n t e r c o n   n e c t e d   p o r e s   i n   t h e   o x i d e .   T h i s   p r o c e s s   b e c o m e s   l e s s   i m p o r t a n t   a t   i n c r e a s i n g   t e m p e r a t u r e s   a s   t h e   o x i d e   s i n t e r s   a n d   t h e   e f f e c t i v e   v o l u m e   a v a i l a b l e   f o r   g a s e o u s   d i f f u s i o n   i s   r e d u c e d .   T h i s   m e c h a n i s m   i s   d i s c u s s e d   b e l o w .   G a s e o u s   D i f f u s i o n   T h r o u g h   O x i d e   P o r e s   B e l o w   1 8 0 0 ~   t h e   e x p e r i m e n t a l l y o b s e r v e d   r a t e   c o n s t a n t s   a r e   m u c h   h i g h e r   t h a n   e x p e c t e d   f o r   a m b i p o l a r   d i f f u s i o n   t h r o u g h   t h e   o x i d e   l a t t i c e .   T h u s ,   o x y g e n   i s   t r a n s p o r t i n g   r a p i d l y   t h r o u g h   t h e   o x i d e ,   s u g g e s t i n g   t h a t   a   s h o r t   c i r c u i t   p a t h   e x i s t s .   T h e   r a t e c o n t r o l l i n g   m e c h a n i s m   b e l o w   1 8 0 0 ~   m u s t   b e   \\x0c', '4 3 4   C o u r t r i g h t   e t   a l .   g a s e o u s   d i f f u s i o n   t h r o u g h   t h e   c o n t i n u o u s   p o r e   n e t w o r k   i n   t h e   o x i d e   l a y e r .   A   g a s d i f f u s i o n   m o d e l   h a s   b e e n   d e v e l o p e d   w h i c h   b a l a n c e s   t h e   c o u n t e r f l o w s   o f   0 2   a n d   C O   t h r o u g h   t h e   o x i d e   l a y e r .   L o c a l l y ,   t h e   C O / C O 2 / O 2   r e a c t i o n s   e s t a b l i s h   e q u i l i b r i u m   e v e r y w h e r e   w i t h i n   t h e   s c a l e .   A t   t h e   t e m p e r a t u r e s   u n d e r   c o n s i d e r a t i o n ,   0 2   a n d   C O   m i x t u r e s   a r e   n o t   s t a b l e   a n d   r e a c t   t o   f o r m   C O 2 .   T h i s   i n c o m p a t i b i l i t y   g i v e s   r i s e   t o   t w o   s e p a r a t e   r e g i o n s   w i t h i n   t h e   p o r o u s   o x i d e .   N e a r   t h e   c a r b i d e ,   t h e   g a s e o u s   s p e c i e s   C O ,   C O 2 ,   a n d   N 2   a r e   p r e s e n t .   I n   t h e   o u t e r   r e g i o n   o f   t h e   s c a l e ,   t h e   g a s e o u s   s p e c i e s   p r e s e n t   a r e   0 2 ,   C O 2 ,   a n d   N 2 .   B e t w e e n   t h e s e   t w o   r e g i o n s ,   a   \" f l a m e   f r o n t \"   e x i s t s   i n   w h i c h   0 2   a n d   C O   r e a c t   t o   f o r m   C O 2 .   T h i s   r e a c t i o n   i s   a   \" s i n k \"   f o r   g a s   m o l e c u l e s ,   w h i c h   r e s u l t s   i n   g a s   d i f f u s i o n   f r o m   b o t h   r e g i o n s   t o w a r d   t h e   f l a m e   f r o n t .   A n   a p p r o x i m a t i o n   o f   t h e   S t e f a n M a x w e l l   e q u a t i o n   w a s   u s e d   t o   c a l c u l   a t e   t h e   c o n c e n t r a t i o n   ( C i )   p r o f i l e s   f o r   e a c h   g a s e o u s   s p e c i e s ,   i ,   a c r o s s   b o t h   r e g i o n s   o f   t h e   p o r o u s   h a f n i a   s c a l e   : 9   j = l   i n   w h i c h   D / e f t   i s   a n   e f f e c t i v e   d i f f u s i o n   c o e f f i c i e n t   o f   i ,   J ;   i s   t h e   f l u x   o f   i   i n   t h e   r e g i o n ,   a n d   x   t h e   d i s t a n c e   f r o m   t h e   c a r b i d e / o x i d e   i n t e r f a c e .   T h e   r e a c t i o n   a t   P a r t i a l   P r e s s u r e   ( A t m . )   0 , 9   0 . 8   N 2   0 . 7   0 . 6   0 . 5   0 . 4   0 . 3   0 . 2   %   ~   0 . 1   C ~   0 . 0   \" \\' \"   0 . 0 0 0   0 . 0 0 2   0 . 0 0 4   0 . 0 0 6   0 . 0 0 8   0 . 0 1 0   D i s t a n c e   f r o m   t h e   C a r b i d e O x i d e   I n t e r f a c e   ( c m )   F i g .   6 .   C a l c u l a t e d   p a r t i a l   p r e s s u r e s   a c r o s s   a   p o r o u s   h a f n i a   s c a l e   a t   1 4 0 0 ~   \\x0c', \"O x i d a t i o n   o f   H a f n i u m   C a r b i d e   4 3 5   2 6   2 4   2 2   2 0   1 8   1 6   P a r a b o l i c   1 4   R a t e   C o n s t a n t   1 2   ( m i c r o n s / s e c   l t 2 )   1 0   8   6   4   2   0   r = 5 0   A   I   I   I   I   \\\\ [   I   I   I   I   0   2   4   6   8   1 0   1 2   1 4   1 6   1 8   2 0   P o r o s i t y   ( % )   F i g .   7 .   T h e   p a r a b o l i c   r a t e   c o n s t a n t   a t   1 4 0 0 ~   a s   a   f u n c t i o n   o f   p o r o s i t y   a n d   a v e r a g e   p o r e   r a d i u s .   t h e   c a r b i d e / o x i d e   i n t e r f a c e   i s :   H f C   +   3 C O 2   =   H f O 2   +   4 C O   ( 7 )   w h i l e   t h e   o v e r a l l   r e a c t i o n   s t o i c h i o m e t r y   i s :   H f C   +   2 0 2   =   H f O 2   +   C O 2   ( 8 )   F i g u r e   6   s h o w s   t h e   p a r t i a l   p r e s s u r e   o f   e a c h   g a s   s p e c i e s   a c r o s s   t h e   s c a l e .   T h e   o v e r a l l   f l u x   o f   o x y g e n   a t o m s   i n w a r d   i s   t w i c e   t h a t   o f   c a r b o n   a t o m s   o u t w a r d ,   a s   r e q u i r e d   b y   E q .   ( 8 ) ,   a n d   a t   e a c h   p o i n t   w i t h i n   t h e   s c a l e ,   a   d i f f u s i o n   e q u a t i o n   c a n   b e   w r i t t e n   f o r   e a c h   g a s   s p e c i e s :   d ( D i e f f d e i l ~   S = 0   ( 9 )   d x   \\\\   R T d x   /   w h e r e   S   i s   a   s o u r c e   o r   s i n k   t e r m .   F o u r   e q u a t i o n s   o f   t h i s   t y p e ,   o n e   e a c h   f o r   0 2 ,   C O ,   C O 2 ,   a n d   N 2 ,   g i v e   r i s e   t o   e i g h t   u n k n o w n s   t h a t   i n c l u d e   t h e   d i f f e r e n t   g a s   p r e s s u r e s   a n d   t h e   s o u r c e   t e r m s .   W h e n   t h e   a v e r a g e   p o r e   r a d i u s   i s   s m a l l ,   d i f f u s i o n   i s   c o n t r o l l e d   p r i m a r i l y   b y   K n u d s e n   d i f f u s i o n ,   a n d   a s   t h e   a v e r a g e   p o r e   r a d i u s   i n c r e a s e s ,   m o l e c u l a r   d i f f u s i o n   b e c o m e s   m o r e   i m p o r t a n t .   T h r o u g h   i t e r a t i v e   s o l u t i o n s   o f   F i c k ' s   d i f f u s i o n   e q u a t i o n s   a n d   t h e   t h e r m o d y n a m i c   e q u i l i b r i u m   e q u a t i o n s ,   i t   w a s   p o s s i b l e   t o   c a l c u l a t e ,   i n   a   s e l f   c o n s i s t e n t   m a n n e r ,   t h e   c o n c e n t r a t i o n   p r o f i l e s   f o r   t h e   d i f f e r e n t   g a s   s p e c i e s   a c r o s s   t h e   o x i d e   l a y e r   a n d   u l t i m a t e l y   p r e d i c t   a   p a r a b o l i c   r a t e   c o n s t a n t   f o r   \\x0c\", '4 3 6   C o u r t r i g h t   e t   a l .   o x i d a t i o n   a s   a   f u n c t i o n   o f   o x i d e   p o r o s i t y ,   a v e r a g e   p o r e   r a d i u s ,   a n d   t e m p e r a   t u r e .   T h e   c a l c u l a t e d   p a r t i a l   p r e s s u r e   o f   C O 2   w a s   f o u n d   t o   p e a k   i n   t h e   m i d d l e   o f   t h e   s c a l e ,   w h i l e   t h e   m a x i m a   i n   t h e   0 2   a n d   t h e   C O   o c c u r   a t   t h e   g a s / o x i d e   a n d   o x i d e / c a r b i d e   i n t e r f a c e s ,   r e s p e c t i v e l y   ( F i g .   6 ) .   T e m p e r a t u r e   w a s   f o u n d   t o   h a v e   v i r t u a l l y   n o   e f f e c t   o n   t h e   c a l c u l a t e d   c u r v e s ,   w h i l e   p o r e   d e n s i t y   a n d   a v e r a g e   p o r e   r a d i u s   h a d   l a r g e   e f f e c t s .   T h e   p a r a b o l i c   r a t e   c o n s t a n t   i s   p r e d i c t e d   t o   v a r y   l i n e a r l y   w i t h   o x y g e n   p a r t i a l   p r e s s u r e .   T y p i c a l   p a r a b o l i c   r a t e   c o n s t a n t s   c a l c u l a t e d   f o r   a   f a m i l y   o f   d i f f e r e n t   p o r e   r a d i i   a t   T =   1 4 0 0 ~   ( 1 6 7 3   K )   a r e   s h o w n   i n   F i g .   7 .   T h e   m o d e l   s u c c e s s f u l l y   p r e d i c t e d   v a l u e s   f o r   t h e   r a t e   c o n s t a n t   t h a t   w e r e   c o n s i s t e n t   w i t h   t h e   e x p e r i   m e n t a l   f i n d i n g s .   F o r   t h e   o b s e r v e d   r a t e   c o n s t a n t   o f   a p p r o x i m a t e l y   8 . 5   # m /   s   1 / 2   a   r a n g e   o f   p o r e   r a d i i   a n d   o x i d e   d e n s i t i e s   w e r e   p r e d i c t e d   t h a t   w e r e   c o n s i s t   e n t   w i t h   t h o s e   a c t u a l l y   o b s e r v e d   i n   t h e   m e t a l l o g r a p h y :   a   p o r o s i t y   o f   a b o u t   5 1 0 %   a n d   p o r e   s i z e s   r a n g i n g   f r o m   0 . 0 1 0 . 1 / ~ m .   S U M M A R Y   A N D   C O N C L U S I O N S   O x i d e   g r o w t h   o n   H f C   a n d   H f C 2 5   T a C   w a s   p a r a b o l i c ,   i n d i c a t i n g   p r o   t e c t i v e   o x i d a t i o n   b e h a v i o r .   A t   t e m p e r a t u r e s   b e l o w   1 6 0 0 ~   t h e   k i n e t i c s   w e r e   d o m i n a t e d   b y   g a s e o u s   d i f f u s i o n   a c r o s s   t h e   o x i d e   l a y e r   v i a   p o r e s   i n   t h e   o x i d e ,   a n d   a   m o d e l   d e s c r i b i n g   t h i s   p r o c e s s   i s   d i s c u s s e d .   A b o v e   1 8 0 0 ~   t h e   k i n e t i c s   w e r e   c o n t r o l l e d   b y   a m b i p o l a r   d i f f u s i o n   o f   o x y g e n   t h r o u g h   t h e   o x i d e   l a y e r   a n d   a u g m e n t e d   b y   g a s e o u s   d i f f u s i o n   i n   p o r e s .   T h e   o x i d a t i o n   o f   h a f n i u m   c a r b i d e s   p r o d u c e s   C O   g a s   w h i c h   c a u s e s   t h e   f o r m a t i o n   o f   p o r e s   t h a t   p r o v i d e   a   p a t h   f o r   b o t h   t h e   o u t w a r d   d i f f u s i o n   o f   g a s e o u s   C O   a n d   t h e   i n w a r d   t r a n s p o r t   o f   o x y g e n   t h r o u g h   t h e   p o r o u s   o x i d e .   T h e   p o r o s i t y   w e a k e n s   t h e   o x i d e   a n d   t h e r e b y   i n c r e a s e s   c r a c k i n g ,   b u t ,   t h e   p o r e s   a l s o   v e n t   t h e   C O   p r e s s u r e   b u i l d u p   a t   t h e   c a r b i d e / o x i d e   i n t e r f a c e ,   a n d   t h i s   p r e v e n t s   c a t a s t r o p h i c   d a m a g e .   T h e   k i n e t i c s   o f   o x i d e   f o r m a t i o n   o n   a   m e t a l   c a r b i d e   a r e   l i m i t e d   b y   g a s   d i f f u s i o n   t h r o u g h   p o r e s   i n   t h e   g r o w i n g   s c a l e   a t   t e m p e r a t u r e s   b e l o w   1 6 0 0 ~   T h u s ,   c a r b i d e   a l l o y   a d d i t i o n s   t h a t   m i g h t   e n h a n c e   t h e   s i n t e r a b i l i t y   o f   t h e   g r o w i n g   o x i d e   l a y e r   a n d   t h e r e b y   r e d u c e   m o l e c u l a r   d i f f u s i o n   c o u l d   s i g n i f i   c a n t l y   i m p r o v e   t h e   o x i d a t i o n   b e h a v i o r   a t   l o w e r   t e m p e r a t u r e s .   O x i d e   c r a c k i n g   a n d   s p a l l a t i o n   w e r e   f o u n d   t o   b e   a   p o t e n t i a l   p r o b l e m   f o r   a l l   H f C b a s e d   s y s t e m s .   T h e   c r a c k s   a p p e a r   t o   b e   r e l a t e d   t o   t h e r m a l l y i n d u c e d   s t r e s s e s .   S i n t e r i n g   a i d s   t h a t   i m p r o v e   t h e   l o w t e m p e r a t u r e   ( <   1 2 0 0 ~   p l a s t i c   i t y   o f   t h e   s c a l e   c o u l d   b e   h e l p f u l   i n   t h i s   r e g a r d .   A d d i t i o n s   o f   T a C   a n d   P r C 2   t o   H f C   d o   n o t   i m p r o v e   t h e   o x i d a t i o n   r e s i s t   a n c e .   H a f n i u m   c a r b i d e   o f   l o w   d e n s i t y   ( ~ 7 5 9 0 % ) ,   o r   m o d i f i e d   w i t h   v o l a t i l e   \\x0c', 'O x i d a t i o n   o f   H a f n i u m   c a r b i d e   4 3 7   s e c o n d p h a s e   a d d i t i o n s ,   o x i d i z e s   a t   g r e a t l y   e n h a n c e d   r a t e s   d u e   t o   t h e   i n c r e a s e d   v o i d   f r a c t i o n   a n d   t h e   i n a b i l i t y   o f   t h e   g r o w i n g   o x i d e   t o   f o r m   a   p r o t e c t i v e   s c a l e .   R E F E R E N C E S   1 .   J .   D .   ( 3 a d d   a n d   E .   B .   E v a n s ,   C o r r o s i o n ,   1 7 ( 9 )   1 0 9   ( 1 9 6 1 ) .   2 .   J .   B .   B e r k o w i t z M a t t u c k ,   J .   E l e c t r o c h e m .   S o c .   1 1 4 ( 1 0 ) ,   p p .   1 0 3 0 1 0 3 3   ( 1 9 6 7 ) .   3 .   R .   F .   V o i t o v i c h   a n d   E .   A .   P u g a c h ,   S o y .   P o w d e r   M e t a l l .   M e t .   C e r a m .   1 2 ( 1 1 ) ,   9 1 6 9 2 1   ( 1 9 7 3 ) .   4 .   K .   M a r n o c h ,   H i g h t e m p e r a t u r e   o x i d a t i o n r e s i s t a n t   h a f n i u m t a n t a l u m   a l l o y s .   J .   M e t .   1 2 2 5   ( N o v .   1 9 6 5 ) .   5 .   A .   K .   K u z n e t s o v ,   P .   A .   T i k h o n o v ,   a n d   M .   V .   K r a v c h i n s k a y a ,   Z h u r n .   N e o r g a n .   K h i m .   2 1 ,   1 3 1 7   ( 1 9 7 6 ) .   6 .   R .   R u h   a n d   V .   A .   P a t e l ,   J .   A m .   C e r a m .   S o c .   5 6 ( 1 1 ) ,   6 0 6   ( 1 9 7 3 ) .   7 .   D .   W .   S t a c y ,   J .   K .   J o h n s t o n e ,   a n d   D .   R .   W i l d e r ,   J .   A m .   C e r a m .   S o c .   5 5 ( 9 ) ,   4 8 2   ( 1 9 7 2 ) .   8 .   W .   W .   S m e l t z e r   a n d   M .   T .   S i m n a d ,   A c t a   M e t a l .   5 ,   3 2 8   ( 1 9 5 7 ) .   9 .   K .   S c h w e r d t f e g e r   a n d   E .   T .   T u r k d o g a n ,   i n   P h y s i c o c h e m i c a l   M e a s u r e m e n t s   i n   M e t a l s   R e s e a r c h ,   V o l .   4 ,   P a r t   I ,   R .   A .   R a p p ,   e d .   ( 1 9 7 0 ) ,   p p .   3 2 1   4 0 7 .   \\x0c']"
},{
  "_id": 167,
  "PDF": "Oxidation of Hafnium Carbide in the Temperature Range 1400” to 2060°C.pdf",
  "Text": "['Oxidation of Hafnium Carbide in the Temperature  Range 1400” to 2060°C   C. Brent Bargeron, Richard C. Benson, A. Norman Je t te , and Terry E. Phillips   Applied Physics Laboratory, The J o h n Hopk ins University, Laurel, Maryland   2 0 7 2 3 4 0 9 9   After hafnium carbide has been oxidized at temperatures in  the range of 1400” to 2060°C, three distinct layers are pres ent in the film cross section: (a) a residual carbide layer with  dissolved oxygen in the lattice, (b) a dense-appearing oxide  interlayer containing carbon, and (c) a porous outer layer  of hafnium oxide. Experimental measurements of  layer  thicknesses and oxygen concentrations are combined with  an extended formulation of moving-boundary diffusion the ory to obtain the diffusion constants of oxygen in each of the  three layers. The results indicate that the oxide interlayer is  a better diffusion barrier for oxygen than either of the other  layers. Based on X-ray microanalysis, X-ray diffraction,  and resistance measurements, the interlayer is an oxygen deficient oxide of hafnium with a carbon impurity. The  interlayer hardness equals that of the residual carbide  layer.   I .   Introduction   H s tance for values o f n from abou t 0.50 to 0.98, a l though , a s  A F N I U M CARB IDE (HfC,) ex is ts as a homogeneous sub  can be seen from a phase d iagram , the end po in ts are a func t ion  of tempera ture .  The HfC, crys ta l is face-cen tered cub ic with  the NaC l s truc ture . Depend ing on the value of x , there can be  large numbers o f vacancies at the C pos i t ions . As x decreases in  value from 0.98 to, 0 .50 , the lattice cons tan t decreases from  4 . 6 4 to abou t 4.60 A . A z ircon ium impurity resu l ts in a slightly  the lattice have sma l ler values. ’L’  larger lattice cons tan t , and crys ta ls with oxygen or n i trogen in  Oxygen con tam ina t ion is difficult to avoid because HfC, can  readily ab so rb large quantities of oxygen in to the la t t ice . A s a  resu l t , it is probably mo re accura te to charac ter ize the material  by  the  formu la HfC,O,.  Cons tan t , Kieffer, and Ettmayer’  showed that at 2000”C , the va lue of y cou ld be a s large a s 0.30  before the lattice would  transform to the HfO, s truc ture . A t  1600°C they indicated the max imum y va lue was 0.25 before  the  transformation  took p lace . Barn ier and ThCvenot’ have  demons tra ted that z ircon ium ca rb ide behaves  the same way.  The HfC,O, characterization was noted ear l ier by Zhe lank in ,  Kutsev, and Ormont’ and by Samsonov and Pade rno .X T h u s ,  hafnium carb ide can d isso lve a fa ir ly large quan t i ty of oxygen  before conver t ing to the ox ide , which ind ica tes tha t the ox ida  tion process for hafn ium carb ide inc ludes nontrivial absorp t ion  and diffusion of oxygen in to the lattice a s preliminary s teps .  Bar t le t t , Wadswor th , and Cutler,’ working in the tempera ture  range of 450” to 580°C , a lso though t th is to be true for z irco nium ca rb ide .  Ox ida t ion tests on hafn ium carb ide performed in the 1960s  and ear l ier are ca tegor ized in handbooks by the phrases “ox ida tion in a i r becomes severe , 1 100-1400”C”~” and “ the ox ida t ion  rate for ZrC and H fC increases l inear ly with t ime , tempera ture ,   1. Sm ia lekLcon t r ihu t ing editor   Manuscript N o . 195854. Received March 2 5 , 1992; approved Oc tuhc r 2 2 , 1992.  Supported hy the Department of the Navy.   and oxygen par t ia l pressure .”” Berkow i tz-Ma t tuck’* wrote that  the “ox ida t ion of H fC , at tempera tures of 1790-2000 K and  oxygen pressures of around 10 torr is . . . linear a n d occurs pref eren t ia l ly a long gra in boundar ies .” Th is latter work was con cerned mainly with samp les prepared by arc hea t ing . Kaufman  and Nesor” ox id ized samp les of “hypereutectic hafn ium car bide” which at the ir work ing tempera tures consisted of HfC  plus exce s s ca rbon . They found the material “exhibited more  rapid ox ida t ion at low tempera tures than at high tempera tures .”  It is unfortunate tha t th is last work is most often cited with the  fact om i t ted that the test material was not hafn ium carb ide at  a l l , but ra ther hypereu tec t ic hafn ium carb ide , which is a hetero geneous m ix ture of hafn ium carb ide and carbon at the test  tempera tures .  In recen t years , a number of new and /or improved techn iques  have been deve loped for mak ing passivating films. Improved  me thods of chem ica l vapor depos i t ion (CVD) are ab le to pro vide c leaner films that are more inert to ex treme env ironmen ts .  Recently, we have ox id ized hafn ium carb ide CVD films which  appear to be protective both above and below the monoc l in ic / te tragonal phase transition tempera ture ( 1700°C) o f HfO,. I 4 . l 5  The initially uniform hafn ium carb ide film was 226 p m thick  accord ing to ca l ibra ted m icroscope measuremen ts . Af ter ox ida tion at high tempera ture in an a tmosphere of 7% O2 and 93%  a rgon , the films were cut to expo se a cross sec t ion . Scann ing  e lec tron microscopy,  light microscopy, X-ray m icroana lys is ,  X-ray  d iffrac t ion ,  and  electrical  resistance measuremen ts  revealed that the ox id ized film cons is ts of three d is t inc t layers  ( see F ig . 1 ) : an ou ter monoc l in ic ox ide layer (HfO,), the resid ual hafn ium carb ide layer with a measurab le oxygen grad ien t ,  and an interlayer wh ich is an ox ide con ta in ing carbon.15 In the  light m icroscope the ou ter layer appears wh i te , the interlayer i s  b lack , and the residual carb ide is metallic gray. The se layers  appeared to be well-bonded to one ano ther as d id the carb ide  layer to the graph i te subs tra te because we d id not observe any  interfacial separa t ion o r o ther s igns of adhesion fa i lure . A s will  be shown below, X-ray m icroana lys is and diffraction clearly  indicated that overa l l the interlayer is an ox ide . Because each  layer in the film is d is t inc t and well-defined  interfaces exist  between layers, oxygen can be expec ted to diffuse at different  ra tes in each ma ter ia l . O n e purpose of this paper is to comb ine  theory and exper imen t to deduce the three diffusion cons tan ts .  Ano ther purpose is to present the exper imen ta l characterization  of the interlayer.  In so lv ing the d iffus iona l prob lem , we ex tend the theory of  Danckwer ts ,” .” which encompasses moving boundar ies and  d iscon t inuous d iffusan t concen tra t ions a t  the  in terfaces ,  to  inc lude three layers. T h e theore t ica l mode l appears to fit the  exper imen ta l c ircums tances qu i te well and an analytical solu tion is found for the overall sy s tem .   11.   Experimental Procedure   The hafn ium carb ide film was prepared us ing chemical vapor  depos i t ion by San Fernando Labora tor ies (now defunc t) . Using  X-ray d iffrac t ion , we identified a face-cep tered cub ic structure  with a lattice cons tan t of 4 . 6 0 ? 0.01 A . If one ignores any   I040   \\x0c', 'April 1993   Oxidution ofHujniurn Curbide in the Temperature Rung e 1400\" to 2060°C   1041   HfO,   Substrate   Fig. 1.   Scanning electron micrograph of a cross section of oxidized  hafnium carbide film showing multilayer structure w i th well-defined  interfaces. No signs of spalling or separation appear between layers.  The compactness and lower porosity of the HfO, AC , (interlayer) are  apparent. Oxidation was for 600 s at I865\"C in an atmosphere of 93%  argon and 7% oxygen.   possible oxygen or n i trogen con tam ina t ion , then accord ing to  da ta presen ted by S torms , \\' this value sugges ts the lattice con  ta ined a subs to ich iome tr ic amoun t of carbon correspond ing to a  formu la o f approx ima te ly HfC,,.,. However,  la t t ice cons tan t  da ta presen ted by Con s tan t , Kieffer, and Ettmaye?  indicate tha t  the presence of e i ther oxygen or n i trogen a lso reduces the lat tice spac ing s im i lar to the reduc t ion caused by carbon vacan c ies . Hafn ium carb ide and hafn ium nitride a re totally m isc ib le  and maintain a face-cen tered cub ic s truc ture for a l l propor t ions  of the two compound s . A l though we never observed n i trogen  dur ing X-ray m icroana lys is of a po l ished cross sec t ion , the  Auger l ines of both n i trogen and oxygen were presen t in the  ana lys is of the or ig ina l film surface . O u r X-ray diffraction da ta  were ob ta ined a t an acu te ang le from the or ig ina l surface ; thus  oxygen and nitrogen con tam ina t ion m igh t have confounded  somewha t  the resu l tan t la t t ice cons tan t ,  ind ica t ing tha t  the  HfC,,, represen ts a lower bound . Wet chem ica l ana lys is of the  film wa s not a t temp ted because of the comp l ica t ion of the need  to separa te the th in layer from a graph i te subs tra te .  We ox id ized hafn ium ca rb ide samp les in a n induc t ion fur nace a s descr ibed prev ious ly . \\'4 .15 T h e inner con ta inmen t tube o f  the furnace wa s made of z ircon ia . T h e spec imen tempera ture  was measured w i th an op t ica l pyrome ter by emp loy ing the  techn ique o f compar ing the samp le b lackbody radiation with a  source of known tempera ture . \\' * Af ter ra is ing the tempera ture to  a prede term ined va lue ,  the a tmosphere  in  the furnace was  sw i tched from pure argon to 9 3% argon p lus 7% oxygen for a  prescribed t ime and then sw i tched back to pu re a rgon . S imu l ta neously, the oven was sw i tched o f f . T h e g a s was flowing across  the spec imen a t a ra te of abou t 2 0 cm / s .  Af ter ox ida t ion , we cu t the films in half to expo se cross sec t ions . The se p ieces were embedded in epoxy and polished us ing  a ser ies of d iamond gr i t down to 0.05 p m . A ca l ibra ted light  m icroscope was emp loyed to measure layer th ickness . Before  in troduc t ion in to the scann ing e lec tron m icroscope (SEM ) , the  samp les were coa ted l igh t ly w i th go ld to avoid charg ing by the  e lec tron beam . We performed the m icroana lys is with a Kevex  w indow less , l i th ium-doped s i l icon de tec tor so tha t carbon and  oxygen X-rays cou ld be co l lec ted . Th ree spec imens (2400 s ,  1400°C; 240 s , 2060°C; and 390 s , 2060°C) were chosen for  ex tended X-ray m icroana lyses . In these spec imen s , the layers  were ana lyzed for a tom ic con ten t a s a func t ion of position such   tha t the ana lys is wa s ob ta ined from a vo lume of o n e to a few  cub ic m icrome ters . T h e spec tra were ana lyzed with sof tware  wh ich accoun ted for the escape peak , sub trac ted the back g round , and in tegra ted the various peaks to ob ta in the number  of coun t s i n e a c h .  In o rde r to exp lo i t and enab le X-ray d iffrac t ion and electrical  res is tance measuremen ts , we cu t o n e spec imen cross section  and po l ished it a t a small ang le re la t ive to the surface so that a  broader  sec t ion of  the  interlayer would be expo sed . Th is  allowed for the irradiation of a par t icu lar layer by a co l l ima ted  X-ray beam , m in im iz ing in terference from the ad jacen t layers.  Additionally, we made a mask of a lum inum foil to minimize  fur ther the X-ray i l lum ina t ion of e i ther the residual carb ide o r  the ou te r ox ide layer. T h e a lum inum foil mask a lso served as a  conven ien t internal s tandard dur ing the ana lys is of the resu l ts .  T h e broader s tr ip of interlayer a lso was advan tageous for mak ing e lec tr ica l con tac t .  We made the res is tance measuremen ts by a t tach ing leads  w i th a Ag pa in t in a l inear four-probe conf igura t ion to the pol ished su r face s o f the ind iv idua l layers. T h e spec imens were  then placed i n a cryos ta t and the res is tance was measured a s i t  was cyc led be tween 10 K and room tempera ture . We de ter m ined hardness numbers us ing s tandard Knoop procedures .  Polished  surfaces  were  a l so  emp loyed  for  the  latter  measuremen ts .   111.   Theory   ,   I ,  X 2 ( 0 ) is   In th i s sec t ion we adap t the two-layer trea tmen t of the mov ing boundary prob lem d u e to Danckwerts\\'\".\\'\\'  to inc lude a third  layer. T h e mode l is dep ic ted in F ig . 2 . Loca t ions in each of the  three layers are g iven by independen t , one-d imens iona l coord i na tes x,,,  x , , x2, respectively. Each of the three coord ina te sys tems is fixed in  its par t icu lar layer as shown . Thu s , dur ing  ox ida t ion the sys tems are in re la t ive motion to o n e another. Th i s  occu r s because each layer has a un ique density. Interface posi t ions a t t ime t are deno ted by X , ( t ) where i = 0, I  2 refers to  coord ina te sys tems . X , ( 0 ) is the initial position of an in terface  in sys tem i . T h e ind ices i = 0 ,  and 2 refer, respectively, to  the ou te r ox ide , the interlayer ox ide , and the residual ca rb ide .  X , ( O ) and X , (O ) a re ind ica ted by do t ted l ines , because they are  no t seen in the cross sec t ion af ter ox ida t ion . We know where  located by measuremen t before ox ida t ion . X : ( t )  X , ( O ) is es t ima ted by ( 1 ) noting tha t all of our pho tographs sug ges t that the interlayer ox ide is dense ( i . e . , without voids or  pores) , (2 ) know ing from m icroana lys is and diffraction  (see  Resu l ts) that it is an ox idetype med ium , and (3) ob ta in ing , by  sub trac t ion , the amoun t of ca rb ide dep le ted . T h u s , ( I ) and (2 )  sugges t tha t we emp loy theoretical dens i t ies of pI = 10.1 g /cm3  and p2 = 12.7 g k m \\'  in the formu laf2p2(X2( t) X 2 ( 0 ) ) =  f,p,(X:(t)  X , ( O ) ) from which we so lve for X T ( t ) X , ( O ) .  (We need th is value to so lve f o r k \\' ; below.) In this express ionf \\' ,  is the we igh t frac t ion of Hf in HfO, andf , is the weight fraction  o f Hf  in H fC , , ;  thus ,  the express ion conserves Hf a tom s .  Imp l ic i t in our scheme is the grow th of the ou te r ox ide at X ( : ( t ) ,  the dep le t ion of the in ter layer ox ide a t X ; ( t ) , the grow th of the  in ter layer ox ide at X : ( t ) ,  and the dep le t ion of the carb ide at  X F ( t ) . T h e reader shou ld no te tha t the ox ida t ion process pro ceed s in an order ly fashion with ca rb ide be ing conver ted to  in ter layer and the interlayer be ing transformed more s low ly in to  the fu l ly ox id ized ox ide . We neg lec t all comp l ica t ions due to  the ou twa rd ( i . e . , in the negative direction in F ig . 2) diffusion  o f ca rbon . We cons ider on ly the inward diffusion of oxygen and  recogn ize the fact that a t ou r working tempera tures , 20-30% of  the ca rbon s i tes in the hafn ium ca rb ide lattice must b e occup ied  by oxygen before the s truc ture transforms to an ox ide . \\'  Because coord ina te sys tems a re no t mov ing re la t ive to their  own layer, the d iffus ion equa t ions can be written simply a s   d c , ( x , , t ) / d t = D , [ a \\' c , ( x , , t ) / i l ~ ; ]   (for i = 0 , 1, 2 )   ( 1 )   \\x0c', '1042   Journal of the American Ceramic Society-Bargeron   et a l .   Vol. 76, N o . 4   Residual  carbide   *   ~   *   I   I   I   I   I   I   I   I x 1   w   I   1  *   I Monoclinic  oxide   \\'   I   I   I   I   I   f   I   I   I   I   I   I   I   Initial  carbide   lnterlayer   Total   interlayer  formed   t   Fig. 2.   Schematic diagram of the hafnium carbide oxidation model.  x,,, x,, and x2 indicate coordinate systems which are fixed in each of the  respective layers. During oxidation these systems move with respect to  one another because of volume changes due to structure transformation  taking place at the interfaces. X \\' s indicate the positions of various  interfaces in the cross section of the oxidized f i lm . Solid interfacial  lines are observable in the cross section. Dashed lines represent origi nal interfaces that no longer exist because of layer depletion.   where C , ( x , , t ) is the concen tra t ion of the d iffusan t (oxygen) in  the ith layer a t point x,, and D , is the correspond ing d iffus ion  cons tan t . Equations  ( I ) with boundary cond i t ions have  the  well-known error function (e r f ) solutions\"   where Cp and B , are cons tan ts de term ined by exper imen t . From  our measuremen ts on the case at hand , the movemen t of the  interfaces appears well-described by the parabolic re la t ions   for i = 0 and 1 . T h e first term on e i ther s ide of the equa t ion rep resents the diffusant flux across the interface whereas the sec ond term takes into accoun t grow th and dep le t ion resu l t ing in  moving boundar ies .  The diffusion cons tan ts are de term ined from the boundary  cond i t ions given by Eq . (4). Us ing the so lu t ions of Eq . ( 1 ) and  the express ions ( 3 ) , the equa t ion a t the interface between layers  0 and 1 can be rewritten a s   , f ( y ) = g(z)   with   exp(y\\') e r f (Y ) I   f ( y ) = I / [ Y   g ( z ) = AJ [z erf (z)] + A?   where   (5 )   (6)   ( 7 )   The solution to these equa t ions is ob ta ined graphically by plot ting f and g  to de term ine where  they a re simultaneously  sa t isf ied .   IV.   Results   The exper imen ta l va lues of the parabolic sca l ing constants k ,  and the concen tra t ions required by the boundary cond i t ions (4)  are presen ted in Table I . The concen tra t ion C(X;(t),t)  is con sistent with t h e resu l ts of Cons tan t , Kieffer, and Ettmdyer\\' at  both T = 2060\" and 1400°C. An illustration of the curves gen era ted by express ions (6) and (7) is given  in F ig . 3 , which  show s a well-defined in tersec t ion o f f a n d g . The determination  of the absc issa of the intersection leads d irec t ly to the diffusion  cons tan ts . Table I1 g ives the values of D , at the two tempera tures . We regard this set of va lues as typical of bulk diffusion  for SeveraLreasons. F irs t , the interfaces between the three lay ers a r e ex treme ly sha rp a n d , second , they are parallel to the  subs tra te and the film surface . Finally, microanalysis never  revealed oxygen concen tra ted in grain boundar ies o r d ispersed  in any way that would sugges t specific pathways through the  film  layers. Even  though m icrographs c lear ly demons tra te  porosity in connec t ion with the ou termos t por t ion of the ou ter  ox ide layer, the pores d o not in genera l appear to be connec ted .  We recogn ize , however, tha t diffusion ac ro s s a s ing le pore will  occu r much  fas ter than  through  the solid  i tse lf . Therefore ,  porosity will cause the diffusion cons tan t to be larger.  F igures 4 , 5 , and 6 present represen ta t ive X-ray microanaly s i s spec tra of the residual ca rb ide , the interlayer, and the ou ter  ox ide layer. Ob se rv ing , in particular, the ratios of the various  peak s , the amoun t of oxygen compared to hafnium increases  substantially go ing from the residual carb ide layer to the inter layer, but d o e s not change much mov ing from the interlayer to  the ou te r ox ide . Th i s is demons tra ted very c lear ly in F ig . 7 ,  where  the da ta were genera ted from microanalysis  spectra  ob ta ined a s a func t ion of position across the cross section of a  film ox id ized for 390 s at 2060°C in 93% argon plus 7% oxy gen . At  the c a r b i d e h t e r l a y e r in terface ,  the [O]/[Hf]  ra t io  increases abrup t ly a lmos t to its va lue in the ou termos t portion  of the ox ide .\" Summar iz ing , the in ter layer appears from the  m icroana lys is results to be an ox ide with carbon in it ra ther than  a ca rb ide with oxygen in i t .  The X-ray d iffrac t ion measuremen ts indicate!  that the HfC  had the usual cub ic s truc ture (a = 4 .60 t 0.01 A) and that the  ou ter ox ide layer had the known monoc l in ic s truc ture . I t was  surpr is ing , cons ider ing the sharply de l inea ted interfaces of the  I ) , that the interlayer a lso had the mono ox id ized film (F ig .  c l in ic la t t ice of  the ou te r ox ide layer. The X-ray diffraction  results are presen ted in Table 111 together with se ts of values   Table I .   Scaling Constants and Concentrations   Tempcrature (\"C)   1400  2060   Temperature (\"c)   1400(2400s)  2060(240s)   2060 ( 3 9 0 s )   k,*   2.08  14.26   C ( x ; ) +   1.59   1.41   1.54   k ,   0 .65  2.76   c(x,: )   1 .33   1.27  1.42   k ;   1.47   13.51   c(x\\';)   1.33   1.27  1.42   K ;   0.82  10.75   c(x: )   1.26  1.25  1.40   k ,   1.02   9 .38   c (x2 )   0.25   0 .33  0 .33   *The units for all the scaling constants are pm / s \" * . The scaling constants k,, k , , and  k, were obtained by fitting a parabola to 3 to 10 measured data points of layer thickness  versus time at a particular temperature. k ; and k\\'; were detcrniined as described in the  text. \\'The units fur all concentrations are g/cm\\'. The concentrations were estimated by  measuring the [O] to [Hf] X-ray ratios and using the value5 o f Constant ef a[.\\' to cali brate the oxygen concentration in the carbide at the interlayer carhide interface.   ~~~   \\x0c', 'April 1993   Oxidation of Hufnium Carb ide in the Temperature Range 1400\" to 2060°C   I043   6.25   0.50   0.75   1 .oo   1 .25   Fig. 3.   Graphs of the functionsf(y) and g(z) showing their intersec tion where Eq. ( 5 ) is satisfied for a film oxidized for 240 s at 2060°C in  93% argon plus 7% oxygen.   Table 11.   Oxygen Diffusion Constants   Layer   Outer oxide   Interlayer oxide   Carbide   Temperature   (\"C)   Diffusion constant  (cm\\'/s)     1400   2060   1400  2060   1400  2060   8  .  1 ~   3.0 X   7 . 9 x 10-9   1 . 1 x l o \\'   2 . 6 x   1.6 x   *A t 2060\"C, the diffusion constants are averages of the results of the analysis of  films oxidized for 240 and 390 s . At 140O0C, the diffusion constants were obtained  horn film oxidized for 2400 s .   from Adam and Rogers,\\'\" Ruh et a l . ,2\\'  and Ge l ler et a1.** Com  par ing the interlayer and ou ter ox ide la t t ice con s tan t s , a and b  are both measurab ly larger for the interlayer, resu l t ing in a ne t  volume increase of abou t 1 .4% . I t is conce ivab le that interstitial  carbon present in the in ter layer is respons ib le for the vo lume  increase . A s seen by compar ing the spec tra in F ig s . 5 and 6 ,  there is more carbon in the in ter layer than in the ou te r ox ide .  T h e cons tan ts of bo th layers a re really no t much d ifferen t from  the literature va lues of presumab ly much purer hafn ia . Because  both the interlayer and the ou te r ox ide have s im i lar s truc tures  and because the amoun t of oxygen change s gradua l ly as one  proceeds from the carb ide i in ter layer in terface through the in ter layer in to the ou te r ox ide , o n e wou ld expec t no v is ib le in terface   Energy (keV )   F ig . 5 .   X-ray spectrum of interlayer after oxidation of original HfC   film at 1865°C for 600 s i n 7% oxygen and 93% argon.   be tween the two layers. S ince th is is no t the c a s e , it seem s  l ike ly that the crys ta l s truc tures a t the ox ida t ion tempera ture  must no t be the same ; o therw ise , the sha rp boundary between  the tw o ma ter ia ls would no t b e expec ted .  T h e resu l ts of the res is tance de term ina t ions a re shown in  F igs . 8 and 9. T h e ou te r ox ide wa s h igh ly insu la t ing (> l o \\'  0.U) and n o results a re shown for tha t layer. T h e magn i tude  and tempera ture dependence of the res is tance of the residual  ca rb ide layer (F ig . 9) and tha t of the in ter layer (F ig . 8 ) are  me ta l l icl ike above 50 K . Below 50 K , however, the res is tance  of both layers beg ins to increase ( the in ter layer mo re than the  residual carb ide) w i th decreas ing tempera ture , sugges t ing that  they m igh t ac tua l ly be more properly descr ibed a s highly  degene ra te sem iconduc tors .  T h e Knoop hardness numbers were 970 and 9 8 6 for the  res idua l ca rb ide and interlayer, respectively, whereas tha t o f the  ou ter layer had a va lue of 297 o r abou t 1/3 of the hardness of the  o the r two . The 113 ra t io is cons is ten t w i th pub l ished hardness  numbers for hafn ium ox ide a n d hafn ium ~ a r b i d e . \\' ~ T h e fac t that  the in ter layer i s much less porous than the ou ter ox ide probab ly  con tr ibu tes to its increased ha rdne s s . It is a l so likely t o be a  ma ter ia l w i th a large numbe r of point defec ts which may a lso  increase i ts hardness . T h e loads used for the Knoop inden ter  were 100 a n d 200 g .   V.   Discussion   T h e d iffus ion cons tan ts in Tab le I1 ind ica te that the interlayer  ox ide is a d iffus ion barr ier fo r oxygen under ou r exper imen ta l  cond i t ions . W i th regard to the diffusion cons tan ts for the ou ter      18   1 6   I Hf   I   I      20   l%  16   I   I   I   I   I   Hf   n   Energy (keV)   Fig. 4.  600 s in 7% oxygen and 93% argon. The gold peak is a result of a coat  X-ray spectrum of residual HfC after oxidation at 1865°C for   ing applied to the specimen to prevent charging in SEM.   Energy (keV)   Fig. 6 .   X-ray spectrum of outer oxide after oxidation of original HfC   film at 1865°C for 600 s i n 7% oxygen and 93% argon .   \\x0c', 'I044   0 4 0   0 35   0 30   42 a a: 0 20  0 0 2 5   2   \\\\   u   c 0 0 15   Y   0 10   0 05   0 00   Journal o f t h e Americun Cerumic Society-Burgeron   et a / .   Vol. 76 , N o . 4   I   I   I   I   I   I   I   o x i d e / l n t e r l a y e r   I In ter la ) ier /carb lde   I   I   ,   0   ,   0 \\'   I   I   01   0   50   100 150 100 250 300 353 400   P o s i t i o n ( p m )   Fig. 7. X-ray spectral ratio of O(K)IHf(M) obtained by analyzing the  spectra as a function of position. To calculate the ratios, the number of  counts in th e hafnium M peak (1.64 keV) and the oxygen K peak were  determined by a Gaussian f i t after processing to account for escape  peaks and after removing the background.   ox ide , we know from th ickness measuremen ts tha t th is layer is  much thicker than it wou ld be if the ox ide were nonporous . T h e  lower density probably con tr ibu ted to a larger d iffus ion con  stant in this layer, and ou r results for the ou ter ox ide shou ld not  be app l ied to fully dense HfO,. We a lso no te tha t at the h igher  tempera ture , the ou ter ox ide layer should possess the te tragona l  s truc ture dur ing the ox ida t ion .  Sm i th , Meszaros , and Amata14 used tubes made of ca lc ia-s ta bilized hafnia to de term ine d i f h s i v i t y from the t ime depen  dence of the permea t ion of oxygen through the tube wa l l . The i r  stated diffusion cons tan ts a re 60-80  t imes  larger than our  results for the ou te r ox ide wh ich , a s has been s ta ted , is not fully  den se . Th i s compar ison indicates tha t the resu l ts of Sm i th e t a / .  likely pertain to diffusion through a c ircu i tous path formed by  gaps between s in tered par t ic les forming the walls of their vessel  and a r e not bulk d iffus ion values for HfO,.  Returning  to the diffusion-barrier na ture o f the  interlayer  ox ide , th is subs tance is composed of on ly three e lemen ts (H f ,  C , and 0) . From our m icrographs , this new material appears to  adhere well to both the residual carb ide and the ou ter ox ide .  Specifically, we have seen no crack ing , spa l l ing , o r separa t ion  a t any of its interfaces with e i ther of the o ther two layers. In  add i t ion , we have not observed cracks or voids within the mate rial i tse lf . Our observa t ions abou t the new material sugges t that  it might be a very usefu l , protective h ightempera ture sub s tance . A pr ior i there d o no t seem to be any reasons why the  material cou ld not be produced as a mono l i th ic protective film.  Barnier and ThCvenot have repor ted work on the syn thes is and   Table 111.   X-ray Diffraction Parameters*   Adam  and Rogcrb\\'\"   Ruh   Ge l le r   Wh i te  ou te r ox ide   Black  interlayer   C t ol.\"   (\\'I ul.\"   u ( 4 )  b ( 4 )   5.1156(5)   5.119   5.12   5.10(1)  5.16(1)   5.14(1)   5.1722(5)   5.169   5 . 1 8   5.19(1)   c ( A )   5.294815)   5.290  99.25   5.25   5.2711)  99.1(1)  136.9   5.27(1)  99.2(1)   (3 (4:g)  V ( A )   99.18(8)   98   138 .30   138.15   137.9   138.8   *A l l X-ray diffraction  results  ind ica te monoc l in ic s t ruc tu re s . Our resiilts fo r the  ou te r o x i d e and inlcrlaycr were de te rn i ined by a weighted lit of 23 and 2X diffraction  l inc s , rc spcc t iv r ly , ob ta ined with a Read came ra using filtered C r rad ia t ion . Po lyc ry s  talline a lum inum Coil placed on the su r face of the target ox ide s w a s emp loyed a s a re f  e rence t o ca l ib ra te the c am e r a . T h e a lum inum which had a slit in it a l so se rved a s a  mask so that da ta were co l lec ted f rom on ly o n e laye r at a t ime .   \\'  O  o  9  \\'  I  100   I   0   200   300   Temperature (K)   ,006   Fig. 8.   Resistance vs temperature of the interlayer (HfO,_,C,). The  units on the ordinate are those of sheet resistance.   ho t-press ing of s ing le-phase ZrC,O, and two-phase ZrC,O,, ZrO, materials.\\'  In a repor t on a ceram ics mee t ing , Car tz states  that dur ing ques t ion ing , ThCvenot reported that carbon appears  to s tab i l ize the cub ic form of ZrO,.*\\' Bar t le t t , Wadswor th , and  Cu t ler also sugges t that ZrO, may be stabilized  in the cub ic  form by sma l l amoun ts of carbon.\\' However, X-ray diffraction  of the interlayer indicates tha t the in ter layer has the monoc l in ic  s truc ture . Thu s , carbon has not s tab i l ized a cub ic s truc ture in  this ca se .  In ou r ana lys is we have largely neglected the role that carbon  diffusion m igh t be playing in ou r sys tem even though ou r sam ples have hafn ium ca rb ide o n a subs tra te of ca rbon . Th i s con figuration i s , of cou r se , a diffusion coup le that cou ld provide a  source for carbon d iffus ion in to the ca rb ide . Our X-ray micro ana lys is resu l ts , however, d o no t reveal a large carbon gradient  at the interface in ques t ion . Although we know of no measure ments of the d iffus ion of carbon in hafn ium ca rb ide , Sar ian and  Cr isc ione have inves t iga ted the d iffus ion of carbon through zir con ium  carb ide  in  the  tempera ture  range  to  2 150°C,2h wh ich corresponds well  to ou r range . Calculating  from their resu l ts , the bulk d iffus ion cons tan ts are 2 x  and 3 X  lo\\' cm 2 / s at 1400\" and 2060\"C, respectively. These  are very small compared with our va lues for oxygen diffusion in  hafn ium ca rb ide in Tab le 11. If it is true that the diffusion of car bon in hafn ium ca rb ide is abou t the same ra te as tha t in zirco nium ca rb ide , then the neg lec t of this process is justified in the  carb ide . On the o ther hand , we recogn ize that the amoun t of  carbon present cou ld easily influence the position of the inter face between the two ox ides .  Comb in ing ou r results above w i th our ear l ier findings and  those of o the r s , a reasonab ly comp le te descr ip t ion of this ox i dation process and resultant ma ter ia ls can be presen ted . Con s tan t , Kieffer, and Ettmayer found that hafn ium carb ide can   from 1350\"   z.007   .ooai   .003 l   I   I   0   100   200   300   Temperature (K)   Fig. 9.   Rehistance vs temperature of the residual HfC layer.     \\x0c', 'April 1993   Oxidcition of Hujniurn Cnrbide in the Teniperuture Runge 1400” fo 2060°C   1045   con ta in abou t 30% oxygen at 2000°C before it conver ts to the  ox ide s t r ~ c t u r e . ~  I t was noted ear l ier tha t the first s tep in the  ox ida t ion of hafn ium ca rb ide is the up take of oxygen in to i ts  s truc ture . Me ta l l ic hafn ium a lso d isso lves oxygen in to its bulk  at e leva ted  T h u s , as oxygen d isso lves in to the  hafn ium ca rb ide la t t ice , i t u l t ima te ly reaches a concen tra t ion  where a s truc tura l pha se transition becomes energe t ica l ly favor ab le . When th is occu r s , there is not enough oxygen to form  fu l ly ox id ized HfO,  and  there  i s still cons iderab le ca rbon  rema in ing in the la t t ice . Therefore , the material tha t is formed  ha s a formu la HfO,  where x is sma l l compared to 2 and y is  probably c loser to 0 than 1 . In o the r wo rd s , the material tha t  forms is a reduced ox ide o f hafn ium with ca rbon in the la t t ice .  Th i s is cons is ten t with the X-ray m icroana lys is and d iffrac t ion  results as well as the res is tance measuremen ts . To unders tand  the cons is tency w i th the res is tance de term ina t ions , reca l l that  hafn ium and titanium a re in the s am e chem ica l g roup and that  rutile TiO,, wh ich is an insulator, becomes a degenera te sem i conduc tor when it is reduced to T iO z , , where x is abou t 0 .2 .  Ho1comb2y observed the in ter layer but made no measuremen ts  to charac ter ize i t . However, he did specu la te in his P h .D . thes is  on the ox ida t ion of hafn ium ca rb ide abou t the na ture of the  interlayer: “A dark black layer is shown to lie be tween the gray  carb ide and the white ox ide . I t is believed to be oxygen-def i cient hafn ia .” Ho lcomb ox id ized ho t-pressed hafn ium ca rb ide  powders wh ich were about 90% dens if ied . Therefore , the pow de r was s in tered but still had relatively large channe l s for the  oxygen to pass through . Th i s probab ly resulted in a less prom i nent interlayer than wou ld resu l t from fully dens if ied ma ter ia l .  The integrity and low porosity o f the interlayer, and its well formed in terfaces with the o the r two layers ind ica te tha t po ten tial ga s produc t ion is no t effec t ive in crea t ing a d isrup t ion at  interfaces o r separa t ion be tween layers. In o the r words , pores  or pocke ts of h igh-pressure ga s d o not seem to pose a prob lem  because of the way the carb ide ox id izes . Th i s i s in s trong con  tras t to the h ightempera ture ox ida t ion of hafn ium d ibor ide .”’  T h e oxidation of the d ibor ide resu l ts in large voids and “stove pipe” s truc tures , both po in t ing to the presence of large quan t i ties of h igh-pressure ga s in the g row ing ox ide . T h e ca rb ide , on  the o ther hand , first absorbs oxygen , then is transformed  compac t oxygen-deficient ox ide . T h e transformation occu r s  prior to full ox ida t ion o f the ca rb ide when there is insufficient  oxygen to ox id ize all of the carbon to CO.  Before summar iz ing o u r resu l ts , we have the following com  men ts o n the monoc l in icke tragona l phase  transition which  occu r s in HfO, at abou t 1700°C. The re seem s to ex is t a belief  with attendant anecdo tes tha t there is an eno rmou s vo lume  change assoc ia ted with th is trans i t ion . (We have heard grea ter  than 10% severa l t imes .) Lynch” d iscusses the vo lume change  and g ives a va lue o f 3 .4% based on measured la t t ice cons tan ts  of the tetragonal phase above 1700°C and exper imen ta l la t t ice  cons tan ts and coeff ic ien ts of expans ion of  the monoc l in ic  phase . In check ing the l i tera ture sources and redo ing the ca lcu la t ion , we come up with essen t ia l ly the s am e vo lume change as  Lynch , tha t i s , abou t 3%. T h e usual anecdo ta l comp la in t is that  the vo lume change cau se s the ox ide to c rumb le or turn to pow de r when coo l ing down from test tempera tures above 1700°C.  We have not encoun tered th is prob lem and no te that Pra ter e t  a l . ,3 2 .2y w h o recently  s tud ied samp les made by ho t-press ing  HfC powders , also d id no t observe th is phenomenon . The re a re  undoub ted ly unknown fac tors ( e . g . , con tam inan ts) at work in  the inves t iga t ions cited anecdotally.   to a   V I .   Summary   We have observed tha t three d is t inc t layers rema in af ter haf n ium ca rb ide is ox id ized at h igh tempera ture : (a) a residual haf nium  ca rb ide  layer,  (b) a dense-appear ing hafn ium ox ide  interlayer con ta in ing ca rbon , and (c ) a porous ou ter layer of  hafn ium ox ide . We have comb ined exper imen ta l measuremen ts  of  layer  th icknesses  and oxygen  concen tra t ions w i th  an  ex tended formulation of the mov ing-boundary d iffus ion theory   of D a n c kw e r t s “ ~ ” to ob ta in the d iffus ion cons tan ts o f oxygen in  each of the three layers. Our ca lcu la t ions ind ica te that the inter layer ox ide is a be t ter d iffus ion barr ier for oxygen than e i ther of  the o the r layers. Our observa t ions of the new material sugges t  tha t  it m igh t be a usefu l , pro tec t ive h ightempera ture sub s tance .” A p r io r i there d o not seem to be any reasons why the  material cou ld not be produced as a mono l i th ic pro tec t ive film.  Our measuremen ts have shown that the in ter layer is an oxy gen-deficient ox ide of hafn ium w i th a ca rbon impurity. I t has a  hardness s im i lar to that o f hafn ium carb ide . As a function of  tempera ture , its res is tance i s sugges t ive of that of a degenera te  sem iconduc tor . 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Interface A w l . , 16, 304-308 (1990).  2YG. R . Holcomb, “The High Temperature Oxidation of Hafnium Carbide”;  Ph.D. Dissertation. Oh io State University, Columbus, OH , 1988.  ”C. B. Bargeron, R . C . Bensen, A . N . Jette, R . W. Newman, and E. L .  Paquette, “Oxidation Morphology and Kinetics of Hafnium Dihoride at High  Temperature,” NASA CP-3097, Part 2 , National Aeronautics and Space Admin istration, Washington, DC, 1990; pp . 545-54.   ‘“2. T. Lynch, “Hafnium Oxide”; pp . 193-216 in H i g h Temperature Ox ide s ,  Part 11. Edited by A . M . Alper. Academic Press , New York, 1970.  ’2J. T. Prater, E. L . Courtwright, G . R . Holcomb, G . R . St. Pierre, and R . A .  Rapp, “Oxidation of Hafnium Carbide and HfC with Additions of Tantalum and  Praseodymium,” NASA CP-3097, Part 1, 1990; pp . 197-209.  ”C . B. Bargeron, R . C . Benson, and A . N . Jette, “High-Temperature Diffu sion of Oxygen in Oxidizing Hafnium Carbide Films,” NASA CP-3054, Part I , National Aeronautics and Space Administration, Washington, DC, 1989; pp.  83-94.   n     \\x0c']"
},{
  "_id": 168,
  "PDF": "Oxidation of HfB2-SiC-Ta4HfC5 ceramic material by a supersonic flow of dissociated air.pdf",
  "Text": "['Contents lists available at ScienceDirect   Journal of the European Ceramic Society   journal homepage: www.elsevier.com/locate/jeurceramsoc   Original Article   Oxidation of HfB2-SiC-Ta4HfC5 ceramic material by a supersonic flow of  dissociated air   Elizaveta P. Simonenko a, *, Nikolay P. Simonenko a, Andrey N. Gordeev b,  Anatoly F. Kolesnikov b, Aleksey V. Chaplygin b, Anton S. Lysenkov c, Ilya A. Nagornov a, d,  Vladimir G. Sevastyanov a, Nikolay T. Kuznetsov a   a Kurnakov Institute of General and Inorganic Chemistry of the Russian Academy of Sciences, Leninsky pr., 31, Moscow, 119991, Russia  b Ishlinsky Institute for Problems in Mechanics of the Russian Academy of Sciences, 101-1 pr. Vernadskogo, Moscow, 119526, Russia  c A.A.Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninskii pr. 49, Moskow, 119334, Russia  d Dmitry Mendeleev University of Chemical Technology of Russia, 9 Miusskaya sq. Moscow, 125047, Russia     A R T I C L E I N F O   A B S T R A C T   Keywords:  Ceramics  UHTC  Refractory carbide  Sol-gel processes  Induction plasmatron   The oxidation of an ultra-high-temperature ceramic material (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 produced by  reactive hot pressing at a temperature of 1800  C (pressure 30 MPa, holding time 30 min, Ar) under long-term  exposure (2000s) to a supersonic flow of dissociated air was studied. It was found that the sample surface  temperature, set during heating and oxidation on a high-frequency induction plasmotron, was significantly lower  than for samples unmodified with super-refractory tantalum-hafnium carbide Ta4HfC5. It was also found that  under similar exposure conditions there was no sharp temperature rise to 2500(cid:0) 2700  C - the temperature did  not exceed 1850  C. Features of the oxidised material surface microstructure were noted, in particular, the  existing gradient in the elemental composition and morphology of the oxide particles forming the surface. It was  found that the main crystalline oxidation product was a complex hafnium-tantalum oxide Hf6Ta2O17, which had  a phase stability up to temperatures of ~2250  C, which set it apart from individual hafnium oxide.     1.  Introduction   Ultra-high-temperature ceramics (UHTC) based on ZrB2(HfB2)-SiC  have a complex of useful properties - the high melting point of components (3000(cid:0) 3040 C - ZrB2 [1], 3220(cid:0) 3380 C - HfB2 [1,2], 2824 C  - SiC (decomposition) [3]) and phase stability over a wide temperature  range, good mechanical properties, relatively high emissivity and high  enough resistance to oxidation, including under exposure to atomic  oxygen [4,5,14,15,6-13]. The high heat conductivity of the main  component that increases further when heated up to 2000 C (134 W⋅  m(cid:0) 1⋅ K(cid:0) 1 for ZrB2 and 143 W⋅ m(cid:0) 1⋅ K(cid:0) 1 for HfB2 [16]) makes it possible  to consider such ceramic materials promising for obtaining sharp edges  of hypersonic vehicles [17-26]. Extensive research on the properties of  ZrB2-SiC and HfB2-SiC ceramics has led to the expansion of their applications, for example, as fuel cells for alternative energetics [27-29]  and in solar energetics [30,31].   The undoubted advantages of such UHTCs are also accompanied by  some very significant negative aspects, such as insufficient strength  characteristics and fracture toughness, and poor thermal shock resistance, which does not allow them to be used when cyclically heated to  high temperatures.  Various approaches are currently being developed to address these  issues:  1) the concept of high entropy alloys is adapted to oxygen-free refractory ceramics [32-38],  2) high-temperature ceramic matrix composites with an antioxidant  and heat-conducting matrix based on MB2-SiC are being developed to a  deeper level [21,39-43], and  3) research is being conducted on the effect of alloying additives of  various nature on the basic characteristics of UHTC [44-46].  It is known [45,47-51] that the introduction of super-refractory  carbides,  such as ZrC/HfC,  into  the  composition of UHTC  can   * Corresponding author.  E-mail addresses: ep_simonenko@mail.ru (E.P. Simonenko), n_simonenko@mail.ru (N.P. Simonenko), koles@ipmnet.ru (A.F. Kolesnikov), alchapl87@gmail.com  (A.V. Chaplygin), toxa55@bk.ru (A.S. Lysenkov), il.nagornov.chem@gmail.com (I.A. Nagornov), vg_sevastyanov@mail.ru (V.G. Sevastyanov), ntkuz@igic.ras.ru  (N.T. Kuznetsov).    https://doi.org/10.1016/j.jeurceramsoc.2020.10.001  Received 6 July 2020; Received in revised form 1 October 2020; Accepted 1 October 2020         \\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            significantly increase the strength of the obtained materials at room  temperature (up to 700(cid:0) 800 MPa). In this case, zirconium and hafnium  carbides are less oxidation resistant, and the available data on the  oxidation of ceramic composites MB2-SiC-MC (M = Zr, Hf) in furnaces in  the atmosphere at temperatures of 1200(cid:0) 1700 C [52-55] indicate a  corresponding decrease in the oxidation resistance and composites as a  whole. However, under conditions of express heating to temperatures  >2000 C with electric heating [56] or using an oxyacetylene torch [52,  57], the authors suggest that ceramic composites containing ZrC are  more stable, due to the formation of a denser layer of ZrO2 on the  surface.  The effect of the introduction of super-refractory tantalum carbide  on the properties of UHTC is much less studied. Thus, there are almost  no data on changes in mechanical characteristics in the literature; there  is only a study [24] that defines the stiffness, elastic modulus and  thermal resistance of the ZrB2-SiC-TaC interface. However, there is a  conflicting report [58-60] on the dependence of ceramics resistance to  oxidation at temperatures of 1200(cid:0) 1600 C on the TaC content - there is  evidence both of its decrease and increase due to the introduction of  tantalum carbide into the material. In an interesting study [59], a higher  oxidation stability of the ZrB2-20 vol%SiC material containing a fairly  high amount of TaC (30 vol%) was established at a temperature of  1500 C, which was explained by the authors in terms of the peculiarities  of distribution of immiscible melts of the vitreous phases of SiO2 and  Ta2O5 in the oxidised layer volume. The complex nature of the oxide  scale and oxidation resistance variation as a function of the oxidation  temperature in a ZrB2-TaSi2 ceramic has been thoroughly described in  [61].  Previously [62], we showed that, when oxidising ceramic materials  (HfB2-30  vol%SiC)-xTa4HfC5  (x = 5, 10, 15  vol%)  containing  super-refractory complex tantalum-hafnium carbide [63-65], under  thermal analysis conditions (20(cid:0) 1400 C, heating speed 20  /min, air  flow rate 250 mL/min), as the content of the carbide component  increased, an increase in mass gain due to oxidation was observed,  indicating a certain decrease in resistance to oxidation. The experiment  performed did not correctly simulate the operating conditions of ceramics, which presumably should be associated with high-speed aerodynamic heating by air flow. This conclusion is confirmed by arc-jet and  ablation tests of ceramic materials based on hafnium and tantalum  carbides carried out in [66,67], the results of which differ significantly  from experiments in static air. We have not found data on the behaviour  of ultra-high-temperature ceramics of the HfB2-SiC-TaC-HfC system in  high-enthalpy air flows in the literature.  In this regard, the purpose of the work reported here was to evaluate  the behaviour of an ultra-high-temperature ceramic material (HfB2-30  vol%SiC)-10 vol%Ta4HfC5 obtained by reactive hot pressing and containing a nanocrystalline complex tantalum-hafnium carbide under  exposure to a supersonic flow of dissociated air.   2.  Experimental procedure   Reagents used: hafnium acetylacetonate [Hf(O2C5H7)4] (99 %) synthesised from HfOCl2⋅8H2O by reacting with acetylacetone C5H8O2  (pure grade) and a 5 % solution of ammonia hydrate NH3⋅H2O; solution  of pentaamyloxytantalum Ta(OC5H11)5 prepared by the reaction of  TaCl5 (extra-pure grade) with amyl alcohol (pure grade) when exposed  to dry ammonia; LBS-1 bakelite varnish (phenol-formaldehyde resin, 1 butanol solution). Preparation of a solution of Ta, Hf-containing precursor was performed by introducing hafnium acetylacetonate into a  tantalum pentaamyloxide butanol solution in the molar ratio n(Ta):n  (Hf) = 4:1 [62,68]. After that, to complete the inter-ligand exchange, the  solution was heated to (cid:0) 80 C and mixed for 30 min.  Synthesis of HfB2-30 vol%SiC composite powder was performed  using the sol-gel technology described in detail in [69], using tetraethoxysilane (TEOS) Si(OC2H5)4 (>99.99 %, EKOS-1 JSC), LBS-1 bakelite  varnish (Karbolit OJSC), formic acid CH2O2 (>99 %, Spektr-Chem LLC),   hafnium diboride (>98 %, particle size 2-3 microns, aggregate size  ~20(cid:0) 60 microns, Tugoplavkie Materialy LLC).  The production of ultra-high temperature ceramic (HfB2-30 vol%  SiC)-10 vol%Ta4HfC5 was performed using the method of reactive hot  pressing of  (HfB2-SiC)@(Ta2O5-HfO2-C)  (mainly with a core-shell  structure) composite powder described in detail in [62], using a hot  press by Thermal Technology Inc. (HP20(cid:0) 3560-20 model) [4-6,70-72]  in graphite moulds with diameter of 15 mm. Heating to a temperature of  1800 C in the argon atmosphere occurred at a rate of 10 C/min, the  holding time at this temperature was 30 min and the applied pressure  was 30 MPa. Boron nitride powder was used as mould lubricant. As a  result, samples were obtained in the form of cylinders with a diameter of  15 mm and a thickness of 4.4-4.5 mm.  During the synthesis of the (HfB2-SiC)@(Ta2O5-HfO2-C) composite  powder, the (Ta2O5-HfO2-C) system is deposited on the surface of HfB2 30 vol%SiC particles,  forming a shell. For this, a powder of the  composition HfB2-30 vol%SiC was dispersed in a solution of phenolformaldehyde resin. Thereafter, a solution of Ta,Hf-containing precursor was introduced dropwise into the reaction system.  The molar ratio of metals and phenolformaldehyde resin was  calculated in such a way that, after carbonisation of the system (dynamic  vacuum, 400 C, 2 h), a full course of the carbothermic synthesis reaction  for Ta4HfC5 was ensured:  2Ta2O5 + HfO2 + 17C = Ta4HfC5 + 12CO.  The complex tantalum-hafnium carbide Ta4HfC5 was selected based  on the promising properties of this particular composition. First of all, it  has  a  record melting  point  [63-65], which  is  important  for  high-temperature ceramics. It was found in [73] that it is for this compound that the highest thermal conductivity is observed in the TaC-HfC  system (34.65 W⋅m-1⋅K-1). In addition, for ceramics of the 4TaC⋅1HfC  composition, a sufficiently high resistance to oxidation is noted, close to  the maximum in the TaC - HfC system [74].  X-ray diffraction analysis was performed on an X-ray diffractometer  Bruker D8 Advance (Bruker Ltd.), CuKα radiation, 0.02  resolution. XRD  pattern, X-ray and full profiles analyses were carried out using the  MATCH! - Phase  Identification  from Powder Diffraction, Version  3.8.0.137 (Crystal Impact, Germany) program, with the use of a COD  reference database integrated into it.  Scanning electron microscopy (SEM) data were obtained with a  triple-beam workstation NVision 40 (Carl Zeiss); the elemental composition of microdomains was determined with an EDX system (Oxford  Instruments).  The behaviour of the ceramics produced under exposure to a supersonic flow of dissociated air was studied on a 100-kilowatt high frequency  induction plasmatron VGU-4 [4,5,8,13,25,68]. Induction  plasmatrons of VGU-series were created in IPMech RAS under the supervision of Dr. M. Yakushin. The configuration of the VGU-4 facility is  shown in fig. S1 and S2, and its main parameters are presented in table  S1. A detailed description of the VGU-4 plasmatron is given in [75]. The  sample was introduced into the flow at the generator anode power  supply (N) of the plasmatron of 30 kW (heat flux q1 was 363 W⋅cm-2),  which was further increased in increments of 10 kW-70 kW (q = 779 W⋅  cm-2). The exposure duration at each stage (N = 30(cid:0) 60 kW) was 2 min.  After reaching N = 70 kW, the sample was exposed at this power until  the end of the experiment; the total exposure time was 33 min 20 s  (2000s). The study used a sound nozzle with an output cross section  diameter of 30 mm. The distance from the nozzle to the sample was 25  mm. The air flow rate was 3.6 g/s and the chamber pressure was 12(cid:0) 14  hPa. To determine the average surface temperature of the sample, a  Mikron M-770S infrared pyrometer was used in the spectral ratio pyrometer mode (temperature range of 1000-3000 C, sighting area  diameter of (cid:0) 5 mm). To record the temperature distribution on the   1 Heat fluxes to a water-cooled copper calorimeter were determined in  separate experiments described in [96].   JournaloftheEuropeanCeramicSociety41(2021)1088-10981089\\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            Fig. 1. The microstructure of the cleaved ceramic (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 (a-c): in the contrast mode by the average atomic number (a, c), according to  the secondary electron detector (b), as well as the X-ray diffraction pattern (d) of the composite powder (HfB2-SiC)@(Ta2O5-HfO2-C) (1) and ceramics (2); the inset  contains EDX-analysis data.   sample surface, a Tandem VS-415U thermal imager was used: thermal  images were recorded at a set value of the spectral radiation coefficient ε  at a wavelength of 0.9 μm equal to 0.65. Then, during the analysis of the  thermal imager data, if necessary, the surface temperature values were   adjusted to real values ε.   Fig. 2. The sample fixation pattern in a water-cooled model with a selected sighting area of the spectral-ratio pyrometer (a), photos of the sample during the  experiment (b,c) and a drawing of the relative position of the sample and the plasma jet (d).   JournaloftheEuropeanCeramicSociety41(2021)1088-10981090\\x0c', '3. Results   The density of the sample produced (HfB2-30 vol%SiC)-10 vol%  Ta4HfC5 was 7.4 ± 0.4 g⋅ cm-3, which is 79.0 ± 3.7 % of the calculated  value obtained by the additive method (the density of HfB2 is assumed to  be 11.2 g⋅ cm-3 [76], SiC - 3.2 g⋅ cm-3 [77] and Ta4HfC5 - 14.17 g⋅ cm-3  [78]).  SEM (Fig. 1a-c) confirmed a fairly uniform distribution of HfB2 grains  (an average size of 1(cid:0) 3 μm), SiC (0.2(cid:0) 0.5 μm in size) and Ta4HfC5 (light  inclusions of 30(cid:0) 60 nm in size), which were formed as a result of the  carbothermic  interaction of  the Ta2O5-HfO2-C  system during  the  consolidation of (HfB2-SiC)@(Ta2O5-HfO2-C) composite powder at a  temperature of 1800 C. EDX analysis within a margin of error of this  method did not detect the presence of oxygen in the produced ceramics.  XRD (Fig. 1d) confirmed a complete synthesis of the Ta4HfC5 carbide  phase: in addition to the typical reflections of HfB2 [79] and SiC [80]  phases of the initial composite powder, the X-ray pattern also shows  reflections of the crystalline complex tantalum-hafnium carbide [78].  No traces of impurity phases (SiO2 [81], HfO2 [82] and HfC [83]) were  detected. The average size of Ta4HfC5 crystallites estimated by the  Scherrer method was 23 ± 2 nm.  In order to study the thermal behaviour of the material produced  under exposure to a supersonic high-enthalpy air flow, a cylindrical  sample with a diameter of 15 mm and a thickness of 4.5 mm was fixed in  a composite copper model using narrow strips of paper based on SiC  filamentous crystals so to prevent contact with the model as much as  possible; the sample fixation pattern corresponded to that in [4-6,71]. In  order to minimise heat transfer to the water-cooled model, the sample  was installed with a 1.5 mm protrusion relative to the front surface  (Fig. 2a). The holder was oriented vertically; the supersonic flow of  dissociated air was directed perpendicular to the sample face surface. A  more detailed diagram of the tool and the mutual orientation of the  sample and the gas flow is shown in Fig. S1-S2 and Fig. 2d.  Introduction of the fixed sample into the supersonic flow was carried  out at the anode power supply of the plasmatron of 30 kW. Photos of the  sample with a model during the experiment, including those showing  the flow phenomena, are shown in Fig. 2b,c. The mode of changing the  power of the anode supply and pressure in the pressure chamber of the  plasmatron, as well as the values of the average surface temperatures  determined using a spectral-ratio pyrometer, are shown in Fig. 3.  As can be seen, at the initial stages of exposure, the sample behaved  similarly to HfB2-SiC ceramic that were subjected to a similar exposure  [4,5,8] (experiments were performed on the same VGU-4 instrument   Fig. 3. The average surface temperature (infrared pyrometer) in the vicinity of  a stagnation point configuration of the sample of (HfB2-30 vol%SiC)-10 vol%  Ta4HfC5, depending on the duration of the experiment and the exposure parameters — the anode power supply N and the pressure in the plasmatron  chamber P.   T T  a  b  l  v e  1  h  e  a  e  r  a  g  e  s  u  r  f  a  c  e  t  e  m  p  e  r  a  t  u  r  e  o  f  s  a  m  p  l  e  s  b  a  s  e  d  o  n  t  h  e  H  f  B  2   S  i  C  s  y  s  t  e  m  d n u  e  r  e  x  p  o  s  u  r  e  t  o  a  s  p u  e  s r  o  n  i  c  fl  o  w  o  f  d  i  s s  o  c  i  a  t  e  d  a  i  r  ,  d  e  p  e  d n  i  n  g  o  n  t  h  e  t  h  e  r  m  a  l  l  o  a  d  a  :  N  -  g  e  n  e  r  a  t  o  r  a  n  o  d  e  p  o  w  e  r  s  p p u  l  y  ,  q   h  e  a  t  fl  u  x  .  E  x a  o p  s  u  r  e e  p  r  a  m  e  t  s r  A  v  e  r  a  g  e  s  u  r  f  a  c  e  t  e  m  p  e  r  a  t  u  r  e  (  p  y  r  o  m  e  t  e  ) r  ,     C  N  ,   k  W  q  ,  W  ⋅  c  m  -  2  (  H  f  B  2   0 3  o v  l  %  S  i  C  )   0 1  o v  l  %  T  a  4  H  f  C  5  ,  o p  r  o  s  i  t  y  5 8 9 3 5 5 5 3 2 3 5 5 8 1 2 3 4 5 2 1 1 1 1 1 1  .  %  H  f  B  2   0 3  o v  l  %  S  i  C  ,  o p  r  o  s  i  t  y  9  %  [  8 6 4 1 1 9 0 4 3 5 3 5 6 7 9 5 1 1 1 1 1  ]  H  f  B  2   0 3  o v  l  %  S  i  C  ,  o p  r  o  s  i  t  y  0 1  .  8  %  [  7 3 5 9 0 9 0 2 0 0 6 3 5 6 7 8 1 1 1 1 1  ]  H  f  B  2   5 6  o v  l  %  S  i  C  ,  o p  r  o  s  i  t  y  4 3  .  5  %  [  8 5 0 0 0 0 1 4 7 2 4 4 5 6 7 6 1 1 1 1 2  ]  (  H  f  B  2   0 3  o v  l  .  %  S  i  C  )   5  o v  l  .  %  Y  3  A  l  5  O  2 1  ,  o p  r  o  s  i  t  y  5  .  5  %  [  4 5 8 0 0 2 6 0 2 9 9 1 4 0 4 5 7 8 3 1 1 1 1 2  ]  0 0 0 0 0 3 4 5 6 7  3 4 8 1 9 6 8 9 9 7 3 4 5 6 7  →  0 5 1 1  → → → → →  3 3 2 4 3 9 1 3 2 4 3 5 6 8 5 1 1 1 1 2  →  5 0 4 1  → → → → →  1 3 4 4 4 6 2 5 1 5 4 6 7 9 5 n 1 1 1 1 2  -  7 4 6 0 1 7 8 1 6 7 5 5 1 1 2 1  → → →  9 7 8 4 8 4 0 3 4 8 2 1 1 1 n  → → →  → → →  0 5 0 4 2 6 7 6 5 1 2 2  →  0 9 5 3 2 2  (  m  i  s  )  b  (  5 2  m  i  n  )  b  (  5 2  m  i  n  )  b  (  m  i  n  0 3  s  )  b  (  ~  m  i  )  b  a  h P  r  e e  s s  u h  r  e  i i  n n  t  h  e  p  l  a a  s  m  a  t  r  o  n  p  r  e  s s  u u  r  e  c  d h n  a  m  o b p  e  r  i  n  a  l l  e  x  p  e  r  i  m  e e  n h  t  s  w  a  s  2 1  (cid:0)  4 1 n  0 P 7 h  a  .  b  T  o  l  d  g  t  i  m  e  t  t  h  e  m  x a  i  m  m  a  o  e  w  e  r  s  p p u  l  y  o  f  t  p  l  a  s  m  a  t  r  o  o  f  k  W  .  E.P. Simonenko et al.                                                                                                                                                                                                                            JournaloftheEuropeanCeramicSociety41(2021)1088-10981091                                                                                                                                                                                                                                                                                                                      \\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            Fig. 4. Thermal images illustrating the temperature distribution over the surface of (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 sample at various stages of exposure to a  supersonic flow of dissociated air and along the sample diameter indicated by a dashed line.   with the parameters specified in Table 1). As the anode power supply of  the plasmatron increased, a gradual increase in the surface temperature  occurred. However, the values of the temperature set on the surface  were significantly lower in this case (Table 1). Thus, the average surface  temperature at an initial power of 30 kW (q = 363 W⋅ cm-2) was set at  1150(cid:0) 1155 C, which  is approximately 250 C  less than the corresponding values observed with a similar exposure of samples to HfB2-30  vol%SiC [5,6] and HfB2-65 vol%SiC [4], and 300(cid:0) 330 C less than for  the sample of HfB2-30 vol%SiC modified by 5 vol% Y3Al5O12 [84].  As N increased to 40 kW, the sample surface temperature (HfB2-30  vol%SiC)-10 vol%Ta4HfC5 increased to ~1240 C, which is even more  different from that for previously tested samples (Table 1).  During the transition to the third and fourth stages of heating (N = 50(cid:0) 60 kW), the average surface temperature increased (at constant N)  at an increasing rate: (cid:0) 10  /min for N = 50 kW and (cid:0) 27  /min for N = 60  kW, which indicates that significant changes in the composition and  microstructure occurred on the surface. The power increase to the  highest value (70 kW, q = 779 W⋅ cm-2) led not only to the expected  gradual increase in the surface temperature by 70(cid:0) 90 C, but also to an  additional increase in the heating rate to ~48  /min. In this case, after  approximately 12 min and before the switching off of the plasmatron,  the surface heating rate was significantly reduced to (cid:0) 5  /min. The  maximum surface  temperature of  the (HfB2-30 vol%SiC)-10 vol%  Ta4HfC5 sample during exposure to a supersonic flow of dissociated air  was 1848 C (for 2000s of the experiment), despite the fact that the  phenomenon of ’temperature jump’ was observed for all samples based  on the HfB2-SiC system [4-6,84], obtained and tested in the same way  (the sample and fixation geometry, the equipment used, heat flux and  pressure, holding time at maximum load) [85]. It was expressed in a  sharp increase in the surface temperature when it reached values of  ~1750(cid:0) 1850 C to 2500(cid:0) 2700 C due to evaporation from the surface of  not only boron oxide, but also silicon oxide, and the appearance of   porous low-heat conducting and highly catalytic hafnium oxide. In this  case, reaching a temperature of (cid:0) 1750(cid:0) 1850 C did not lead to rapid  and sharp heating.  In addition, it should be noted that, even with lower heat flux, the  temperature of the oxidised surface of the sample containing 10 vol%  Ta4HfC5 was several hundred degrees lower than for samples of HfB2 SiC [4,5,8] and, especially, (HfB2-30 vol% SiC)-5 vol% Y3Al5O12 [84].  The distribution of the temperature field on the surface of samples  during their oxidation was studied using a thermal imager (Fig. 4). As  can be seen from the thermal images and the temperature profile along  the diameter, there was a temperature gradient [4-6,84] for the material  (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 that is typical for this type of material under exposure to a supersonic air flow — the highest surface  temperature was concentrated in the central region and gradually  decreased towards the edge of the sample. As the anode power supply of  the plasmatron and the exposure duration increased, ΔT increased from  ~100 C (N = 30 kW, q = 363 W⋅ cm-2, t = 1 min) to 230 C (N = 70 kW,  q = 779 W⋅ cm-2, t = 33 min 20 s).  As a result of the sample exposure to a supersonic flow of dissociated  air, a mass increase of 5.9 % ((cid:0) 0.2 g⋅ cm-2) was observed after 2000s,  and the oxidised sample layer thickness increased by (cid:0) 0.02 mm (0.4 %).  A study of the oxidised surface microstructure using SEM makes  possible to distinguish three significantly different regions: central  ((cid:0) 6(cid:0) 8 mm in diameter, Fig. 5), intermediate (Fig. 6) and peripheral  (Fig. 7). As can be seen from photomicrographs, the surface morphology  was completely different from the typical surface of HfB2-SiC materials  after oxidation: instead of a porous ceramic crust, there was a hierarchically organised structure based on flat plates oriented towards the  surface at an angle close to 90  . The phase contrast data make it possible  to conclude that these plates had a similar elemental composition. In  deeper layers, where the plates adhered to the surface, relatively thin  interlayers of a less absorbing substance, probably silicate glass with an   JournaloftheEuropeanCeramicSociety41(2021)1088-10981092\\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            Fig. 5. Microstructure and elemental composition of the central region of the surface oxidised by supersonic air flow for the sample of (HfB2-30 vol%SiC)-10 vol%  Ta4HfC5: in contrast mode by the average atomic number (a,c,e,f), according to the secondary electron detector (b,d); accelerating voltage of 1 kV (a-d) and 20 kV  (e,f).   admixture of tantalum oxide, can be seen at the borders. This assumption is confirmed by photomicrographs made with a high accelerating  voltage of 20 kV (Fig. 5e,f), which show even more contrastingly the  oxidised surface microstructure, which was a conglomerate of plates  ~300(cid:0) 500 nm thick and 1-5 μm long, probably bonded together by  interlayers of silicate glass. The general appearance of the material  resembled the oxidised surface (1627 C, 100 min, in stagnant air) of the  ZrB2-20 vol%SiC-20 vol%TaC sample shown in [58], in which vertically  oriented flat particles also began to appear under the silicate glass layer.  EDX analysis of the oxidised surface in the central region of the sample  showed that the molar ratio n(Hf):n(Ta) was 3.06, which within the  error range of the method corresponds to the composition of the individual complex oxide Hf6Ta2O17 [86,87]. The amount of silicon on the  surface was so small that its content was not recorded.  In the intermediate region between the centre and the periphery of  the sample (Fig. 6), the oxidised surface microstructure represented hilly   formations of a highly dispersed Hf, Ta-containing phase (Fig. 6d)  covered with silicate glass (Fig. 6a-c). In this case, the surface was  dominated by silicon oxide: according to the EDX data of the entire  micrograph field, the molar ratio n(Si):n(Hf) was 1.5. The ratio n(Hf):n  (Ta) = 1.21 can be explained by the sufficiently high content of Ta2O5 in  the vitreous layer.  On the sample periphery after oxidation, bulges of 10(cid:0) 50 μm were  formed over the ceramic surface enriched with hafnium (Fig. 7, EDX),  the number of which decreased towards the edge (Fig. 7a). The  elemental composition of the ceramic layer showed that the ratio n(Hf):n  (Ta) was slightly less than that for Hf6Ta2O17 and was 2.6; silicon oxide  was not detected according to EDX data - Fig. 7d. Local analysis of  protruding formations that were close to spherical in shape showed that  they contained mainly silicon dioxide (the ratio of n(Si):n(Hf) was quite  high, at (cid:0) 7.8). They were probably bubbles of silicate glass squeezed out  of deeper  layers through the pores of hafnium-tantalum oxide by   JournaloftheEuropeanCeramicSociety41(2021)1088-10981093\\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            Fig. 6. Microstructure and elemental composition of the intermediate region of the surface oxidised by supersonic air flow for the sample of (HfB2-30 vol%SiC)-10  vol%Ta4HfC5: in contrast mode by the average atomic number; accelerating voltage of 1 kV (a-c) and 20 kV (d).   excessive pressure of gaseous oxidation products. In this case, it should  also be noted that, in the composition of such bubbles, the amount of  tantalum oxide exceeded the amount of hafnium oxide (n(Hf):n(Ta) = 0.6), in contrast to the surface over which they protruded, enriched with  HfO2.  The qualitative distribution of the main elements (Hf, Si) over the  oxidised surface on the sample periphery is demonstrated by mapping in  Fig. 8, which clearly shows that hafnium oxide was mainly concentrated  between rounded formations containing silicon.  X-ray diffraction analysis of the oxidised surface of the (HfB2-30 vol  %SiC)-10 vol%Ta4HfC5 sample showed (Fig. 9) that its phase composition differed significantly from that for UHTC doped with TaC [58-60].  Thus, in this work, as a result of oxidation of ZrB2-20 vol%SiC-(5÷30)vol  %TaC materials in the temperature range of 1500(cid:0) 1700 C, the formation of either a mixture of ZrO2, ZrSiO4, SiO2, Ta2O5 phases [59,60], or a  mixture of tetragonal and monoclinic ZrO2 phases, was observed [58].  In this work, the highest intensity was found in the reflections of the  Hf6Ta2O17 orthorhombic phase [86,87], which, in contrast to HfO2, had  phase stability over a wide range of temperatures — up to peritectic  decay at a temperature of 2244 ± 37 C [88]. In addition to the main  phase of Hf6Ta2O17, the X-ray pattern shows low-intensity reflections of  monoclinic and orthorhombic Ta2O5 [89,90], probably crystallised from  silicate glass.   4. Discussion   Summarising the data obtained directly during the test, and as a  result of the sample examination after exposure to a supersonic flow of  dissociated air, a number of peculiarities of the oxidation process can be  identified.  First of all, it should be noted that the temperature values that were  set when heating the (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 ceramic were   reduced in comparison with unmodified HfB2-30 vol%SiC material -  Table 1. In this case, the mode of exposure to the samples was similar -  the  same plasmatron with well-reproducible exposure parameters  (anode power supply and pressure in the plasmatron chamber, the mode  of their change). It can be assumed that this was due to a certain increase  in the heat conductivity of ceramic when doped with nanodispersed  complex carbide Ta4HfC5.  In addition to the ’temperature jump’, usually observed during  prolonged (>15(cid:0) 20 min) exposure to high-enthalpy air flow, was not  observed in this case. This phenomenon is associated with evaporation  of a low-catalytic borosilicate glass melt from the surface and with the  corresponding exposure of the porous framework of highly-catalytic and  low-heat-conducting HfO2, which, of course, has a radically lower  vapour pressure at these temperatures than the evaporation products of  SiO2 and, especially, B2O3 [91-95]. That is, a kind of thermal barrier  layer is created that concentrates heat on the surface and prevents its  removal to the volume of the ceramics.  It should be noted that the gradual increase in the sample surface  temperature observed after 4 min of heating at a constant anode power  supply N ≥ 50 kW indirectly indicates that the chemical composition of  the surface changed in this case, which was confirmed by SEM with EDX  (Figs. 5-7). However, the process of SiO2 evaporation from the surface is  probably much slower.  Reducing the partial pressure of SiO vapour over silicate glass is  possible not only if there is a lower surface temperature, but also if the  glass contains tantalum oxide, which has more than two orders of  magnitude lower vapour pressure at a temperature of 1800(cid:0) 1850  C  (Fig. S3). The EDX data of glass bubbles in the peripheral region of the  sample (Fig. 7) confirm that tantalum oxide formed during the oxidation  of Ta4HfC5 (at a temperature lower than HfB2) is consumed not only for  the formation of the Hf6Ta2O17 phase, but also enters the composition of  silicate glass. Here, the n(Ta):n(Hf) ratio was 1.8, while for Hf6Ta2O17   JournaloftheEuropeanCeramicSociety41(2021)1088-10981094\\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            Fig. 7. Microstructure and elemental composition of the peripheral region of the surface oxidised by supersonic air flow for the sample of (HfB2-30 vol%SiC)-10 vol%  Ta4HfC5: in contrast mode by the average atomic number (a-e), according to the secondary electron detector (a); accelerating voltage of 1 kV.   Fig. 8. Distribution of Hf and Si elements on the oxidised surface of the sample of (HfB2-30 vol%SiC)-10 vol%Ta4HfC5, peripheral region.    the n(Ta):n(Hf) ratio was 0.33. The ratio of the intensity of XRD peaks of  the oxidised surface (Fig. 9) indicates that the amount of crystalline  Ta2O5 was low, i.e. the main part of the Ta2O5 was probably part of the  glassy phase melt.  Therefore, the combined factors of lower surface temperature and  reduced SiO2 vapour pressure over the SiO2-Ta2O5 melt (possibly in the  form of two immiscible liquids [59,61]) led to a decrease in the evaporation rate from the glassy layer surface and, accordingly, to the fact  that the phenomenon of the usually observed ’temperature jump’ from  1800(cid:0) 1850  C to ~2500(cid:0) 2700 C under prolonged exposure to the  supersonic flow of dissociated air, in this case, did not occur. Most likely,   the surface temperature excursion should also occur during the oxidation  of  ultra-high-temperature materials  doped with  complex  tantalum-hafnium carbides, but this will require an increase in the  duration of exposure and heat flux.   5. Conclusion   In this work, we studied the oxidation of an ultra-high-temperature  ceramic material (HfB2-30 vol%SiC)-10 vol%Ta4HfC5 produced by  reactive hot pressing at a temperature of 1800 C (pressure 30 MPa,  holding time 30 min, Ar) under exposure to a supersonic flow of   JournaloftheEuropeanCeramicSociety41(2021)1088-10981095\\x0c', 'E.P. Simonenko et al.                                                                                                                                                                                                                            SiC-(5÷30)vol%TaC  in the temperature range of 1500(cid:0) 1700 C  in  stagnant air [58-60], and not monoclinic HfO2 with a very small  admixture of Ta2O5 as we had observed in the oxidation of (HfB2-30 vol  %SiC)-10 vol%Ta4HfC5 material in the thermal analysis mode in the  temperature range of 20(cid:0) 1400 C in [62], but an orthorhombic complex  oxide Hf6Ta2O17, which had a phase stability up to temperatures of  (cid:0) 2250 C, which sets it apart from individual hafnium oxide.  In general, the research has shown the great potential of the ultra high-temperature ceramic (HfB2-30 vol%SiC)-10 vol%Ta4HfC5, which,  despite long-term harsh exposure to a supersonic flow of dissociated air,  acquired an average surface temperature of <1900 C, on the surface of  which a phase-stable complex oxide Hf6Ta2O17 was formed, in a wide  temperature range as a result of oxidation.  The data obtained indicate the necessity to continue systematic  research of the oxidative behaviour of ceramics based on the HfB2-SiC  system, modified with super-refractory carbides.   Declaration of Competing Interest   The authors declare that they have no known competing financial  interests or personal relationships that could have appeared to influence  the work reported in this paper.   Acknowledgements   This work was supported by the Russian Foundation for Basic  Research (project no. 20-03-00502). The SEM and X-ray phase analysis  measurements were performed using shared experimental facilities  supported  by  IGIC RAS  state  assignment  (no.  01201353364).  <GN3>Ishlinsky Institute for Problems in Mechanics RAS state assignment<GN3?>. Experiments on the HF-plasmatron VGU-4 were partially  carried out within the framework of state assignment of Ishlinsky  Institute  for  Problems  in Mechanics  RAS  (no.  АААА-А20 120011690135-5).   Appendix A.  Supplementary data   Supplementary material related to this article can be found, in the  online version, at doi:https://doi.org/10.1016/j.jeurceramsoc.2020.10  .001.   References   [4]  [1] X.B. Wang, D.C. Tian, L.L. Wang, The electronic structure and chemical stability of  the AlB2-type transition-metal diborides, J. Phys. Condens. Matter 6 (1994)  10185-10192.  [2] H. Bittermann, P. Rogl, Critical assessment and thermodynamic calculation of the  ternary system boron-hafnium-titanium (B-Hf-Ti), J. Phase Equilibria. 18 (1997)  24-47, https://doi.org/10.1007/BF02646757.  [3] K. Korniyenko, Boron - Carbon - Silicon, in: Refract. Met. Syst. Phase Diagrams,  Crystallogr. Thermodyn. Data. Landolt-b¨ornstein, New Ser. IV/11E1, Springer,  2009, pp. 499-534, https://doi.org/10.1007/978-3-540-88053-0_21.  E.P. Simonenko, N.P. Simonenko, A.N. Gordeev, A.F. Kolenikov, A.S. Lysenkov, I.  A. Nagornov, V.G. Sevastyanov, N.T. 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Combined data  from  the spectral-ratio pyrometer and  thermal  imager showed that the surface temperature change during heating  differed significantly from that observed for UHTC of the HfB2-SiC system [4,5,8], including the modified 5 vol% Y3Al5O12 [84]. In this case,  for a sample containing 10 vol% of nanocrystalline Ta4HfC5, lower  surface temperatures were observed (by 250(cid:0) 450 C), which may have  been due to the sufficiently high heat conductivity of the produced  material. It was found that, when the surface temperature reached  ~1750(cid:0) 1850 C, even at the maximum anode power supply of the  plasmatron of 70 kW (the heat flux measured  in relation to the  water-cooled copper calorimeter was 779 W⋅ cm-2), there was no  ’temperature  jump’, as previously noted  for samples of a similar  composition subjected to a similar exposure to a supersonic flow of  dissociated air. Perhaps the reason for this phenomenon is associated  with the decrease in the evaporation activity of silicon dioxide due to the  introduction of Ta2O5 into the glass melt or due to an increase in its  viscosity.  It is noted that, for the sample of (HfB2-30 vol%SiC)-10 vol%Ta4HfC5  when oxidised by a high-enthalpy air flow, there was no stabilisation of  the surface temperature at a fixed value of q = 779 W⋅ cm-2; after the  12th min of heating, the temperature increase rate was ~5 C/min. The  highest value of the surface-averaged temperature according to the  spectral-ratio pyrometer was 1848 C (for 2000s of the experiment).  As a result of oxidation, a mass gain of 5.9 % was recorded, which is  also a significant difference from the behaviour of UHTC in [4-6,84],  where there was a mass loss from 1.7 (for a sample of (HfB2-30 vol%  SiC)-5 vol% Y3Al5O12 [84]) to 8.6 % (for a high-porosity sample of  HfB2-65 vol%SiC [4]).  The study of the microstructure and elemental composition of the  oxidised surface showed that the central part of the sample differed  significantly from the periphery, which corresponds to the data on the  existence of a temperature gradient (ΔT to 230 C). In the central region,  a hierarchical structure was formed based on Hf6Ta2O17 plates oriented  perpendicular to the sample surface; the presence of SiO2 impurity was  not detected using EDX analysis. The content of silicon oxide increased  towards the edge of the sample, partly due to the formation of silicate  glass bubbles containing Ta2O5 on the surface of the ceramic layer of  hafnium-tantalum oxides.  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  "_id": 169,
  "PDF": "Oxidation of HfB2-SiC-Ta4HfC5ceramic material by a supersonic flow of dissociated air.pdf",
  "Text": "['Journal Pre-proof  Oxidation of HfB2 -SiC-Ta4HfC5 ceramic material by a supersonic flow of dissociated air  Elizaveta P. Simonenko, Nikolay P. Simonenko, Andrey N. Gordeev, Anatoly F. Kolesnikov, Aleksey V. Chaplygin, Anton S. Lysenkov, Ilya A. Nagornov, Vladimir G. Sevastyanov, Nikolay T. Kuznetsov  PII:  DOI:  S0955-2219(20)30813-X  https://doi.org/10.1016/j.jeurceramsoc.2020.10.001  Reference:  JECS 13629  To appear in:  Journal of  the European Ceramic Society  Received Date:  6 July 2020  Revised Date:  1 October 2020  Accepted Date:  1 October 2020  Please cite this ar ticle as: { doi: https://doi.org/  This is a PDF file of an ar ticle that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the ar ticle. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal per tain.  © 2020 Published by Elsevier.  \\x0c', 'Russia, 125047, e-mail: il.nagornov.chem@gmail.com    Leninskii pr. 49, Moskow, Russia, 119334, e-mail: toxa55@bk.ru   dDmitry Mendeleev University of Chemical Technology of Russia, 9 Miusskaya sq. Moscow ,   Oxidation of HfB2-SiC-Ta4HfC5 ceramic material by a supersonic flow of dissociated air   Elizaveta   P.   Simonenkoa,*,   Nikolay   P.   Simonenkoa,   Andrey   N.   Gordeevb,   Anatoly F. Kolesnikovb, Aleksey V. Chaplyginb, Anton S. Lysenkovc, Ilya A. Nagornova,d,   Vladimir G. Sevastyanova, Nikolay T. Kuznetsova   aKurnakov Institute of General and Inorganic Chemistry of the Russian Academy of Sciences,   Leninsky   pr.,   31,   Moscow,   Russia,   119991,   e-mail:   ep_simonenko@mail.ru,   n_simonenko@mail.ru, vg_sevastyanov@mail.ru, ntkuz@igic.ras.ru   bIshlinsky Institute for Problems in Mechanics of the Russian Academy of Sciences, 101-1 pr.   Vernadskogo, Moscow, Russia, 119526, e-mail: koles@ipmnet.ru, alchapl87@gmail.com   cA.A.Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences,   Aleksey V. Chaplygin, ORCID https://orcid.org/0000-0001-9606-6095   Nikolay P. Simonenko, ORCID https://orcid.org/0000-0002-4209-6034   Elizaveta P. Simonenko, ORCID http://orcid.org/0000-0001-8112-1821   Journal Pre-proof  UHTC (HfB2-30 vol% SiC)-10 vol%Ta4HfC5 was obtained by reactive hot pressing   Its behaviour when exposed to a very-high-speed flow of dissociated air was studied   UHTC HfB2-30 vol% SiC   were determined   Abstract   A significant difference was shown in comparison with the behaviour of unmodified   Highlights   \\uf0b7   \\uf0b7   \\uf0b7   \\uf0b7   The features of the oxidised surface microstructure, elemental and phase composition   The oxidation of an ultra-high-temperature ceramic material   (HfB2-30vol%SiC) 10vol%Ta4HfC5 produced by reactive hot pressing at a temperature of 1800\\uf0b0С (pressure   30 MPa, holding time 30 min, Ar) under long-term exposure (2000 s) to a supersonic flow of   dissociated air was studied. It was found that the sample surface temperature, set during heating   and oxidation on a high-frequency induction plasmotron, was significantly lower than for   1               \\x0c', 'samples unmodified with super-refractory tantalum-hafnium carbide Ta4HfC5. It was also   found that under similar exposure conditions there was no sharp temperature rise to 2500 2700\\uf0b0С - the temperature did not exceed 1850\\uf0b0С. Features of the oxidised material surface   microstructure were noted, in particular, the existing gradient in the elemental composition and   morphology of the oxide particles forming the surface. It was found that the main crystalline   oxidation product was a complex hafnium-tantalum oxide Hf6Ta2O17, which had a phase   stability up to temperatures of ~2250\\uf0b0С, which set it apart from individual hafnium oxide.   Keywords: Ceramics; UHTC; Refractory   carbide; Sol-gel processes;   Induction   plasmatron   useful properties - the high melting point of components (3000-3040\\uf0b0C - ZrB2 [1], 3220 3380\\uf0b0C - HfB2 [1,2], 2824\\uf0b0C - SiC (decomposition) [3]) and phase stability over a wide   temperature range, good mechanical properties, relatively high emissivity and high enough   resistance to oxidation, including under exposure to atomic oxygen [4,5,14,15,6-13]. The high   heat conductivity of the main component that increases further when heated up to 2000\\uf0b0С   (134 W\\uf0d7m-1\\uf0d7K-1 for ZrB2 and 143 W\\uf0d7m-1\\uf0d7K-1 for HfB2 [16]) makes it possible to consider such   Ultra-high-temperature ceramics (UHTC) based on ZrB2(HfB2)-SiC have a complex of   1.   Introduction   Journal Pre-proof  Various approaches are currently being developed to address these issues:    1) the concept of high entropy alloys is adapted to oxygen-free refractory ceramics [32-  The undoubted advantages of such UHTCs are also accompanied by some very   ceramic materials promising for obtaining sharp edges of hypersonic vehicles [17-26].   Extensive research on the properties of ZrB2-SiC and HfB2-SiC ceramics has led to the   expansion of their applications, for example, as fuel cells for alternative energetics [27-29] and   significant negative aspects, such as insufficient strength characteristics and fracture toughness,   and poor thermal shock resistance, which does not allow them to be used when cyclically heated   in solar energetics [30,31].   to high temperatures.    38],    2) high-temperature ceramic matrix composites with an antioxidant and heat-conducting   matrix based on MB2-SiC are being developed to a deeper level [21,39-43], and   3) research is being conducted on the effect of alloying additives of various nature on   the basic characteristics of UHTC [44-46].   2       \\x0c', 'It is known [45,47-51] that the introduction of super-refractory carbides, such as   ZrC/HfC, into the composition of UHTC can significantly increase the strength of the obtained   materials at room temperature (up to 700-800 MPa). In this case, zirconium and hafnium   carbides are less oxidation resistant, and the available data on the oxidation of ceramic   composites MB2-SiC-MC (M=Zr, Hf) in furnaces in the atmosphere at temperatures of 1200 1700\\uf0b0С [52-55] indicate a corresponding decrease in the oxidation resistance and composites   as a whole. However, under conditions of express heating to temperatures >2000\\uf0b0C with   electric heating [56] or using an oxyacetylene torch [52,57], the authors suggest that ceramic   composites containing ZrC are more stable, due to the formation of a denser layer of ZrO2 on   the surface.   UHTC is much less studied. Thus, there are almost no data on changes in mechanical   report [58-60] on the dependence of ceramics resistance to oxidation at temperatures of 1200  1600\\uf0b0С on the TaC content - there is evidence both of its decrease and increase due to the   (30vol%) was established at a temperature of 1500\\uf0b0С, which was explained by the authors in   introduction of tantalum carbide into the material. In an interesting study [59], a higher   oxidation stability of the ZrB2-20vol%SiC material containing a fairly high amount of TaC   resistance  variation as a function of the oxidation temperature in a ZrB2-TaSi2 ceramic has   characteristics in the literature; there is only a study [24] that defines the stiffness, elastic   modulus and thermal resistance of the ZrB2-SiC-TaC interface. However, there is a conflicting   The effect of the introduction of super-refractory tantalum carbide on the properties of   Journal Pre-proof  Previously [62], we showed that, when oxidising ceramic materials (HfB2-30vol%SiC) terms of the peculiarities of distribution of immiscible melts of the vitreous phases of SiO2 and   Ta2O5 in the oxidised layer volume. The complex nature of the oxide scale and oxidation   presumably should be associated with high-speed aerodynamic heating by air flow. This   conclusion is confirmed by arc-jet and ablation tests of ceramic materials based on hafnium and   tantalum carbides carried out   in [66,67],   the results of which differ significantly from   experiments in  static air. We have not found data on the behaviour of ultra-high-temperature   ceramics of the HfB2-SiC-TaC-HfC system in high-enthalpy air flows in the literature.   3   250 ml/min), as the content of the carbide component increased, an increase in mass gain due   to oxidation was observed, indicating a certain decrease in resistance to oxidation. The   experiment performed did not correctly simulate the operating conditions of ceramics, which   xTa4HfC5 (x=5, 10, 15vol%) containing super-refractory complex tantalum-hafnium carbide   [63-65], under thermal analysis conditions (20-1400\\uf0b0С, heating speed 20\\uf0b0/min, air flow rate   been thoroughly described in [61].   \\x0c', 'In this regard, the purpose of the work reported here was to evaluate the behaviour of   an ultra-high-temperature ceramic material (HfB2-30vol%SiC)-10vol%Ta4HfC5 obtained by   reactive hot pressing and containing a nanocrystalline complex tantalum-hafnium carbide under   exposure to a supersonic flow of dissociated air.   2. Experimental Procedure   Reagents used: hafnium acetylacetonate   [Hf(O2C5H7)4]   (99%)   synthesised   from   HfOCl2·8H2O by reacting with acetylacetone C5H8O2 (pure grade) and a 5% solution of   ammonia hydrate NH3·H2O; solution of pentaamyloxytantalum Ta(OC5H11)5 prepared by the   n(Ta):n(Hf)=4:1 [62,68]. After that, to complete the inter-ligand exchange, the solution was   EKOS-1 JSC), LBS-1 bakelite varnish (Karbolit OJSC), formic acid СН2О2 (>99%, Spektr Chem LLC), hafnium diboride   (>98%, particle size 2-3 microns, aggregate size ~20 technology described in detail in [69], using tetraethoxysilane (TEOS) Si(OC2H5)4 (>99.99%,   reaction of TaCl5 (extra-pure grade) with amyl alcohol (pure grade) when exposed to dry   ammonia; LBS-1 bakelite varnish (phenol-formaldehyde resin, 1-butanol solution). Preparation   introducing hafnium   solution   in   the molar   ratio   Synthesis of HfB2-30vol%SiC composite powder was performed using the sol-gel   a   into   tantalum pentaamyloxide butanol   solution of Ta, Hf-containing precursor was performed by   Journal Pre-proof  the (HfB2-SiC)@(Ta2O5-HfO2-C) composite powder,   During   the synthesis of   The production of ultra-high temperature ceramic (HfB2-30vol%SiC)-10vol%Ta4HfC5   was performed using the method of reactive hot pressing of (HfB2-SiC)@(Ta2O5-HfO2-C)   (mainly with a core-shell structure) composite powder described in detail in [62], using a hot   (Ta2O5-HfO2-C) system is deposited on the surface of HfB2-30vol%SiC particles, forming a   shell. For this, a powder of the composition HfB2-30vol%SiC was dispersed in a solution of   phenolformaldehyde resin. Thereafter, a solution of Ta,Hf-containing precursor was introduced   dropwise into the reaction system.   4   in the form of cylinders with a diameter of 15 mm and a thickness of 4.4-4.5 mm.   press by Thermal Technology Inc. (HP20-3560-20 model) [4-6,70-72] in graphite moulds with    diameter of 15 mm. Heating to a temperature of 1800\\uf0b0С in the argon atmosphere occurred at a   rate of 10°C/min, the holding time at this temperature was 30 min and the applied pressure was   30 MPa. Boron nitride powder was used as mould lubricant. As a result, samples were obtained   of a   acetylacetonate   heated to ~80\\uf0b0С and mixed for 30 min.    60 microns, Tugoplavkie Materialy LLC).   the     \\x0c', 'database integrated into it.   NVision 40 (Carl Zeiss); the elemental composition of microdomains was determined with an   EDX system (Oxford Instruments).   dissociated air was studied on a 100-kilowatt high-frequency induction plasmatron VGU-4   The molar ratio of metals and phenolformaldehyde resin was calculated in such a way   that, after carbonisation of the system (dynamic vacuum, 400°C, 2 h), a full course of the   carbothermic synthesis reaction for Ta4HfC5 was ensured:   2Ta2O5 + HfO2 + 17C = Ta4HfC5 + 12CO.   The complex tantalum-hafnium carbide Ta4HfC5 was selected based on the promising   properties of this particular composition. First of all, it has a record melting point [63-65],   which is important for high-temperature ceramics. It was found in [73] that it is for this   compound   that   the highest   thermal conductivity   is observed   in   the TaC-HfC system   (34.65 W\\uf0d7m-1 K-1). In addition, for ceramics of the 4TaC\\uf0d71HfC composition, a sufficiently high   resistance to oxidation is noted, close to the maximum in the TaC - HfC system [74].   Advance (Bruker Ltd.), CuK\\uf061 radiation, 0.02\\uf0b0 resolution. XRD pattern, X-ray and full profiles   analyses were carried out using the MATCH! - Phase Identification from Powder Diffraction,   Version 3.8.0.137 (Crystal Impact, Germany) program, with the use of a COD reference   The behaviour of the ceramics produced under exposure to a supersonic flow of   Scanning electron microscopy (SEM) data were obtained with a triple-beam workstation   X-ray diffraction analysis was performed on an X-ray diffractometer Bruker D8   Journal Pre-proof  4 plasmatron is given in [75]. The sample was introduced into the flow at the generator anode   power supply (N) of the plasmatron of 30 kW (heat flux q1 was 363 W\\uf0d7cm-2), which was further   increased in increments of 10 kW to 70 kW (q=779 W\\uf0d7cm-2). The exposure duration at each   [4,5,8,13,25,68]. Induction plasmatrons of VGU-series were created in IPMech RAS under the   supervision of Dr. M. Yakushin. The configuration of the VGU-4 facility is shown in figures S1   nozzle to the sample was 25 mm. The air flow rate was 3.6 g/s and the chamber pressure was   12-14 hPa. To determine the average surface temperature of the sample, a Mikron M-770S   infrared pyrometer was used in the spectral ratio pyrometer mode (temperature range of 1000-  1Heat fluxs to a water-cooled copper calorimeter were determined in separate experiments described   in [96].   5   stage (N=30-60 kW) was 2 min. After reaching N=70 kW, the sample was exposed at this power   until the end of the experiment; the total exposure time was 33 min 20 s (2000 s). The study   used a sound nozzle with an output cross section diameter of 30 mm. The distance from the   and S2, and its main parameters are presented in table S1. A detailed description of the VGU                                                   \\x0c', '3000\\uf0b0C, sighting area diameter of ~5 mm). To record the temperature distribution on the sample   surface, a Tandem VS-415U thermal imager was used: thermal images were recorded at a set   value of the spectral radiation coefficient \\uf065 at a wavelength of 0.9 µm equal to 0.65. Then,   during the analysis of the thermal imager data, if necessary, the surface temperature values were   adjusted to real values \\uf065.   ResultsThe density of the sample produced (HfB2-30vol%SiC)-10vol%Ta4HfC5 was   7.4±0.4 g\\uf0d7cm-3, which is 79.0±3.7% of the calculated value obtained by the additive method   (the density of HfB2 is assumed to be 11.2 g\\uf0d7cm-3 [76], SiC - 3.2 g\\uf0d7cm-3 [77] and Ta4HfC5 -   14.17 g\\uf0d7cm-3 [78]).    composite powder, the X-ray pattern also shows reflections of the crystalline complex tantalum hafnium carbide [78]. No traces of impurity phases (SiO2 [81], HfO2 [82] and HfC [83]) were   detected. The average size of Ta4HfC5 crystallites estimated by the Scherrer method was   consolidation of (HfB2-SiC)@(Ta2O5-HfO2-C) composite powder at a temperature of 1800\\uf0b0С.   the produced ceramics. XRD (Fig. 1d) confirmed a complete synthesis of the Ta4HfC5 carbide   phase: in addition to the typical reflections of HfB2 [79] and SiC [80] phases of the initial   of 1-3 µm), SiC (0.2-0.5 µm in size) and Ta4HfC5 (light inclusions of 30-60 nm in size), which   were formed as a result of the carbothermic interaction of the Ta2O5-HfO2-C system during the   SEM (Fig. 1a-c) confirmed a fairly uniform distribution of HfB2 grains (an average size   Journal Pre-proof  In order to study the thermal behaviour of the material produced under exposure to a   surface (Fig. 2a). The holder was oriented vertically; the supersonic flow of dissociated air was   directed perpendicular to the sample face surface. A more detailed diagram of the tool and the   thickness of 4.5 mm was fixed in a composite copper model using narrow strips of paper based   on SiC filamentous crystals so to prevent contact with the model as much as possible; the   sample fixation pattern corresponded to that in [4-6,71]. In order to minimise heat transfer to   the water-cooled model, the sample was installed with a 1.5 mm protrusion relative to the front   mutual orientation of the sample and the gas flow is shown in Fig. S1-S2 and Fig. 2d.    Introduction of the fixed sample into the supersonic flow was carried out at the anode   power supply of the plasmatron of 30 kW. Photos of the sample with a model during the   experiment, including those showing the flow phenomena, are shown in Fig. 2b,c. The mode   of changing the power of the anode supply and pressure in the pressure chamber of the   6   EDX analysis within a margin of error of this method did not detect the presence of oxygen in   23±2 nm.    supersonic high-enthalpy air flow, a cylindrical sample with a diameter of 15 mm and a     \\x0c', \"plasmatron, as well as the values of the average surface temperatures determined using a   spectral-ratio pyrometer, are shown in Fig. 3.    As can be seen, at the initial stages of exposure, the sample behaved similarly to HfB2 SiC ceramic that were subjected to a similar exposure [4,5,8] (experiments were performed on   the same VGU-4 instrument with the parameters specified in table.1). As the anode power   supply of the plasmatron increased, a gradual increase in the surface temperature occurred.   However, the values of the temperature set on the surface were significantly lower in this case   (Table 1). Thus, the average surface temperature at an initial power of 30 kW (q=363 W\\uf0d7cm-2)   [84].    tested samples (Table 1).   (HfB2-30vol%SiC) surface temperature increased (at constant N) at an increasing rate: ~10\\uf0b0/min for N=50 kW and   ~27\\uf0b0/min for N=60 kW, which indicates that significant changes in the composition and   but also   increase   in   the heating rate   to ~48\\uf0b0/min. In   this case, after   approximately 12 min and before the switching off of the plasmatron, the surface heating rate   was set at 1150-1155\\uf0b0С, which is approximately 250\\uf0b0С less than the corresponding values   observed with a similar exposure of samples to HfB2-30vol%SiC [5,6] and HfB2-65vol%SiC   10vol%Ta4HfC5 increased to ~1240\\uf0b0С, which is even more different from that for previously   [4], and 300-330\\uf0b0С less than for the sample of HfB2-30vol%SiC modified by 5vol% Y3Al5O12   microstructure occurred on the surface. The power increase to the highest value (70 kW, q=779   W\\uf0d7cm-2) led not only to the expected gradual increase in the surface temperature by 70-90\\uf0b0С,   During the transition to the third and fourth stages of heating (N=50-60 kW), the average   As N   increased   to 40 kW,   temperature   the sample surface   Journal Pre-proof  to an additional   jump' was observed for all samples based on the HfB2-SiC system [4-6,84], obtained and tested   in the same way (the sample and fixation geometry, the equipment used, heat flux and pressure,   holding time at maximum load) [85]. It was expressed in a sharp increase in the surface   was significantly reduced to ~5\\uf0b0/min. The maximum surface temperature of the (HfB2 30vol%SiC)-10vol%Ta4HfC5 sample during exposure to a supersonic flow of dissociated air   temperature when it reached values of ~1750-1850\\uf0b0С to 2500-2700\\uf0b0С due to evaporation from   the surface of not only boron oxide, but also silicon oxide, and the appearance of porous low heat conducting and highly catalytic hafnium oxide. In this case, reaching a temperature of   ~1750-1850\\uf0b0С did not lead to rapid and sharp heating.   In addition, it should be noted that, even with lower heat flux, the temperature of the   oxidised surface of the sample containing 10vol%Ta4HfC5 was several hundred degrees lower   7   was 1848\\uf0b0С (for 2000 s of the experiment), despite the fact that the phenomenon of 'temperature   \\x0c\", 'than for samples of HfB2-SiC [4,5,8] and, especially, (HfB2-30 vol% SiC)-5 vol% Y3Al5O12   [84].   The distribution of the temperature field on the surface of samples during their oxidation   was studied using a thermal imager (Fig. 4). As can be seen from the thermal images and the   temperature profile along the diameter, there was a temperature gradient [4-6,84] for the   material (HfB2-30vol%SiC)-10vol%Ta4HfC5 that is typical for this type of material under   exposure to a supersonic air flow — the highest surface temperature was concentrated in the   central region and gradually decreased towards the edge of the sample. As the anode power   increase of 5.9% (~0.2 g\\uf0d7cm-2) was observed after 2000 s, and the oxidised sample layer   thickness increased by ~0.02 mm (0.4%).    oriented towards the surface at an angle close to 90\\uf0b0. The phase contrast data make it possible   to conclude that these plates had a similar elemental composition. In deeper layers, where the   supply of the plasmatron and the exposure duration increased, \\uf044T increased from ~100\\uf0b0С   (N=30 kW, q=363 W\\uf0d7cm-2, t=1 min) to 230\\uf0b0С (N=70 kW, q=779 W\\uf0d7cm-2, t=33 min 20 s).   three significantly different regions: central (~6-8 mm in diameter, Fig. 5), intermediate (Fig.   6) and peripheral (Fig. 7). As can be seen from photomicrographs, the surface morphology was   plates adhered to the surface, relatively thin interlayers of a less absorbing substance, probably   silicate glass with an admixture of tantalum oxide, can be seen at the borders. This assumption   completely different from the typical surface of HfB2-SiC materials after oxidation: instead of   a porous ceramic crust, there was a hierarchically organised structure based on flat plates   A study of the oxidised surface microstructure using SEM makes  possible to distinguish   As a result of the sample exposure to a supersonic flow of dissociated air, a mass   Journal Pre-proof  layer. EDX analysis of the oxidised surface in the central region of the sample showed that the   molar ratio n(Hf):n(Ta) was 3.06, which within the error range of the method corresponds to   the composition of the individual complex oxide Hf6Ta2O17 [86,87]. The amount of silicon on   the surface was so small that its content was not recorded.    In the intermediate region between the centre and the periphery of the sample (Fig. 6),   the oxidised surface microstructure represented hilly formations of a highly dispersed Hf, Ta 8   conglomerate of plates ~300-500 nm thick and 1 to 5 µm long, probably bonded together by   interlayers of silicate glass. The general appearance of the material resembled the oxidised   is confirmed by photomicrographs made with a high accelerating voltage of 20 kV (Fig. 5e,f),   which show even more contrastingly   the oxidised surface microstructure, which was a   surface (1627\\uf0b0С, 100 min, in stagnant air) of the ZrB2-20vol%SiC-20vol%TaC sample shown   in [58], in which vertically oriented flat particles also began to appear under the silicate glass     \\x0c', 'amount of tantalum oxide exceeded the amount of hafnium oxide (n(Hf):n(Ta)=0.6), in contrast   to the surface over which they protruded, enriched with HfO2.   containing phase (Fig. 6d) covered with silicate glass (Fig. 6a-c). In this case, the surface was   dominated by silicon oxide: according to the EDX data of the entire micrograph field, the molar   ratio n(Si):n(Hf) was 1.5. The ratio n(Hf):n(Ta)=1.21 can be explained by the sufficiently high   content of Ta2O5 in the vitreous layer.   On the sample periphery after oxidation, bulges of 10-50 µm were formed over the   ceramic surface enriched with hafnium (Fig. 7, EDX), the number of which decreased towards   the edge (Fig. 7a). The elemental composition of the ceramic layer showed that the ratio   n(Hf):n(Ta) was slightly less than that for Hf6Ta2O17 and was 2.6; silicon oxide was not detected   according to EDX data - Fig. 7d. Local analysis of protruding formations that were close to   spherical in shape showed that they contained mainly silicon dioxide (the ratio of n(Si):n(Hf)   was quite high, at ~7.8). They were probably bubbles of silicate glass squeezed out of deeper   layers through the pores of hafnium-tantalum oxide by excessive pressure of gaseous oxidation   products. In this case, it should also be noted that, in the composition of such bubbles, the   The qualitative distribution of the main elements (Hf, Si) over the oxidised surface on   analysis of   the oxidised   X-ray diffraction   Journal Pre-proof  surface of   the   (HfB2-30vol%SiC) monoclinic ZrO2 phases, was observed [58]. In this work, the highest intensity was found in the   reflections of the Hf6Ta2O17 orthorhombic phase [86,87], which, in contrast to HfO2, had phase   intensity reflections of monoclinic and orthorhombic Ta2O5 [89,90], probably crystallised from   10vol%Ta4HfC5 sample showed (Fig. 9) that its phase composition differed significantly from   that for UHTC doped with TaC [58-60]. Thus, in this work, as a result of oxidation of ZrB2 of either a mixture of ZrO2, ZrSiO4, SiO2, Ta2O5 phases [59,60], or a mixture of tetragonal and   stability over a wide range of temperatures — up to peritectic decay at a temperature of   2244±37\\uf0b0С [88]. In addition to the main phase of Hf6Ta2O17, the X-ray pattern shows low the sample periphery is demonstrated by mapping in Fig. 8, which clearly shows that hafnium   oxide was mainly concentrated between rounded formations containing silicon.    20vol%SiC-(5÷30)vol%TaC materials in the temperature range of 1500-1700\\uf0b0С, the formation   silicate glass.   3. Discussion    Summarising the data obtained directly during the test, and as a result of the sample   examination after exposure to a supersonic flow of dissociated air, a number of peculiarities of   the  oxidation process can be identified.   9     \\x0c', \"First of all, it should be noted that the temperature values that were set when heating the    (HfB2-30vol%SiC)-10vol%Ta4HfC5 ceramic were reduced in comparison with unmodified   HfB2-30vol%SiC material - Table 1. In this case, the mode of exposure to the samples was   similar - the same plasmatron with well-reproducible exposure parameters (anode power   supply and pressure in the plasmatron chamber, the mode of their change). It can be assumed   that this was due to a certain increase in the heat conductivity of ceramic when doped with   nanodispersed complex carbide Ta4HfC5.    In addition to the 'temperature jump', usually observed during prolonged (>15-20 min)   exposure to high-enthalpy air flow, was not observed in this case.  This phenomenon is   volume of the ceramics.   much slower.    after 4 minutes of heating at a constant anode power supply N \\uf0b3 50 kW indirectly indicates that   the  chemical composition of the surface  changed in this case, which was confirmed by SEM   with EDX (Fig. 5-7). However, the process of SiO2 evaporation from the surface is probably   associated with evaporation of a low-catalytic borosilicate glass melt from the surface and with   the corresponding exposure of   low-heat conducting HfO2, which, of course, has a radically lower vapour pressure at these temperatures   than the evaporation products of SiO2 and, especially, B2O3 [91-95]. That is, a kind of thermal   barrier layer is created that concentrates heat on the surface and prevents its removal to the   It should be noted  that the gradual increase in the sample surface temperature observed   the porous framework of highly-catalytic and   Journal Pre-proof  Reducing the partial pressure of SiO vapour over silicate glass is possible not only if   of silicate glass. Here, the n(Ta):n(Hf) ratio was 1.8, while for Hf6Ta2O17 the n(Ta):n(Hf) ratio   was 0.33. The ratio of the intensity of  XRD peaks of the oxidised surface (Fig. 9) indicates that   that tantalum oxide formed during the oxidation of Ta4HfC5 (at a temperature lower than HfB2)   is consumed not only for the formation of the Hf6Ta2O17 phase, but also enters the composition   the amount of crystalline Ta2O5 was low, i.e. the main part of the Ta2O5 was probably part of   more than two orders of magnitude lower vapour pressure at a temperature of 1800-1850\\uf0b0С   (Fig. S2). The EDX data of glass bubbles in the peripheral region of the sample (Fig. 7) confirm   there is a lower surface temperature, but also if the glass contains tantalum oxide, which has   the glassy phase melt.    Therefore, the combined factors of lower surface temperature and reduced SiO2 vapour   pressure over the SiO2-Ta2O5 melt (possibly in the form of two immiscible liquids [59, 61]) led   to a decrease in the evaporation rate from the  glassy layer surface and, accordingly, to the fact   that the phenomenon of the usually observed 'temperature jump' from 1800-1850°C to ~2500 2700°C under prolonged exposure to the supersonic flow of dissociated air, in this case, did not   10   \\x0c\", \"to 779 W\\uf0d7cm-2.   surface temperature change during heating differed significantly from that observed for UHTC   occur. Most likely, the surface temperature excursion should also occur during the oxidation of   ultra-high-temperature materials doped with complex tantalum-hafnium carbides, but this will   require an increase in the duration of exposure and heat flux.   4. Conclusion   In this work, we studied the oxidation of an ultra-high-temperature ceramic material   (HfB2-30vol%SiC)-10vol%Ta4HfC5 produced by reactive hot pressing at a temperature of   1800\\uf0b0С (pressure 30 MPa, holding time 30 min, Ar) under exposure to a supersonic flow of   dissociated air. The sample with a relative density of 81.5% was exposed to a high-enthalpy jet   for 2000 s, while there was a gradual increase in the anode power supply of the plasmatron from   30 to 70 kW and, accordingly, of the heat flux to the water-cooled copper calorimeter from 363   of the HfB2-SiC system [4,5,8], including the modified 5 vol% Y3Al5O12 [84]. In this case, for   a sample containing 10vol% of nanocrystalline Ta4HfC5, lower surface temperatures were   1850\\uf0b0C, even at the maximum anode power supply of the plasmatron of 70 kW (the heat flux   measured in relation to the water-cooled copper calorimeter was 779 W\\uf0d7cm-2), there was no   Combined data from the spectral-ratio pyrometer and thermal imager showed that the   Journal Pre-proof  It is noted that, for the sample of (HfB2-30vol%SiC)-10vol%Ta4HfC5 when oxidised by   As a result of oxidation, a mass gain of 5.9% was recorded, which is also a significant   a high-enthalpy air flow, there was no stabilisation of the surface temperature at a fixed value   of q=779 W\\uf0d7cm-2; after the 12th min of heating, the temperature increase rate was ~5\\uf0b0C/min.   'temperature jump', as previously noted for samples of a similar composition subjected to a   similar exposure to a supersonic flow of dissociated air. Perhaps the reason for this phenomenon   The highest value of the surface-averaged temperature according to the spectral-ratio pyrometer   observed (by 250-450\\uf0b0С), which may have been due to the sufficiently high heat conductivity   of the produced material. It was found that, when the surface temperature reached ~1750 is associated with the decrease in the evaporation activity of silicon dioxide due to the   introduction of Ta2O5 into the glass melt or due to an increase in its viscosity.   was 1848\\uf0b0С (for 2000 s of the experiment).    difference from the behaviour of UHTC in [4-6,84], where there was a mass loss from 1.7 (for   a sample of (HfB2-30 vol% SiC)-5 vol% Y3Al5O12 [84]) to 8.6% (for a high-porosity sample of   HfB2-65vol%SiC [4]).   The study of the microstructure and elemental composition of the oxidised surface   showed that the central part of the sample differed significantly from the periphery, which   11     \\x0c\", 'temperature range of 20-1400\\uf0b0C in [62], but an orthorhombic complex oxide Hf6Ta2O17, which   had a phase stability up to temperatures of ~2250\\uf0b0С, which sets it apart from individual hafnium   oxide.   corresponds to the data on the existence of a temperature gradient (\\uf044Т to 230\\uf0b0С). In the central   region, a hierarchical structure was formed based on Hf6Ta2O17 plates oriented perpendicular   to the sample surface; the presence of SiO2 impurity was not detected using EDX analysis. The   content of silicon oxide increased towards the edge of the sample, partly due to the formation   of silicate glass bubbles containing Ta2O5 on the surface of the ceramic layer of hafnium tantalum oxides.   XRD of the surface showed that the main crystalline product of oxidation was not   monoclinic HfO2, as we had noted earlier for materials based on HfB2-SiC [4-6,84], not a   mixture of ZrO2, ZrSiO4, SiO2 and Ta2O5 phases as observed in the oxidation of UHTC ZrB2 20vol%SiC-(5÷30)vol%TaC in the temperature range of 1500-1700\\uf0b0С in stagnant air [58-60],   supersonic flow of dissociated air, acquired an average surface temperature of <1900\\uf0b0С, on the   surface of which a phase-stable complex oxide Hf6Ta2O17 was formed, in a wide temperature   and not monoclinic HfO2 with a very small admixture of Ta2O5 as we had observed in the   oxidation of (HfB2-30vol%SiC)-10vol%Ta4HfC5 material in the thermal analysis mode in the   In general, the research has shown the great potential of the ultra-high-temperature   Journal Pre-proof  Acknowledgements   state assignment   IGIC RAS   This work was supported by the Russian Foundation for Basic Research (project no. 20 The data obtained indicate the necessity to continue systematic research of the oxidative   experimental   facilities   supported by   (no. 01201353364).   Experiments on the HF-plasmatron VGU-4 were partially carried out within the framework of   state assignment of Ishlinsky Institute for Problems in Mechanics RAS no. АААА-А20 behaviour of ceramics based on the HfB2-SiC system, modified with super-refractory carbides.   ceramic (HfB2-30vol%SiC)-10vol%Ta4HfC5, which, despite long-term harsh exposure to a   range as a result of oxidation.    03-00502). The SEM and X-ray phase analysis measurements were performed using shared   120011690135-5.   Declaration of competing interest   The authors declare that they have no known competing financial interests or personal   relationships that could have appeared to influence the work reported in this paper.   12         \\x0c', 'References   [1]   X.B. Wang, D.C. Tian, L.L. Wang, The electronic structure and chemical stability of the   AlB2-type transition-metal diborides, J. Phys. Condens. Matter. 6 (1994) 10185-10192.   [2]   H. Bittermann, P. 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Kuznetsov,   Behavior of HfB2-SiC (10, 15, and 20 vol %) ceramic materials in high-enthalpy air   flows,   Russ.   J.   Inorg.   Chem.   61   (2016)   1203-1218.   https://doi.org/10.1134/S003602361610017X.   Journal Pre-proof  22       \\x0c', 'Captions to Figures   Fig. 1. The microstructure of the cleaved ceramic (HfB2-30vol%SiC)-10vol%Ta4HfC5 (a-c): in   the contrast mode by the average atomic number (a, c), according to the secondary electron   detector (b), as well as the X-ray diffraction pattern (d) of the composite powder (HfB2 SiC)@(Ta2O5-HfO2-C) (1) and ceramics (2); the inset contains EDX-analysis data   Journal Pre-proof  the relative position of the sample and the plasma jet (d).   Fig. 2. The sample fixation pattern in a water-cooled model with a selected sighting area of the   spectral-ratio pyrometer (a), photos of the sample during the experiment (b,c) and a drawing of   23       \\x0c', 'Fig. 3. The average surface temperature (infrared pyrometer) in the vicinity of a stagnation   point configuration of the sample of (HfB2-30vol%SiC)-10vol%Ta4HfC5, depending on the   duration of the experiment and the exposure parameters — the anode power supply N and the   pressure in the plasmatron chamber P,   Journal Pre-proof  dissociated air and along the sample diameter indicated by a dashed line.   Fig. 4. Thermal images illustrating the temperature distribution over the surface of (HfB2 30vol%SiC)-10vol%Ta4HfC5 sample at various stages of exposure to a supersonic flow of   24     \\x0c', 'by the average atomic number (a,c,e,f), according to the secondary electron detector (b,d);   accelerating voltage of 1 kV (a-d) and 20 kV (e,f).    Fig. 5. Microstructure and elemental composition of the central region of the surface oxidised   by supersonic air flow for the sample of (HfB2-30vol%SiC)-10vol%Ta4HfC5: in contrast mode   Journal Pre-proof  25     \\x0c', 'Journal Pre-proof  Fig. 6. Microstructure and elemental composition of the intermediate region of the surface   oxidised by supersonic air flow for the sample of (HfB2-30vol%SiC)-10vol%Ta4HfC5: in   contrast mode by the average atomic number; accelerating voltage of 1 kV (a-c) and 20 kV (d).    26     \\x0c', 'contrast mode by the average atomic number (a-e), according to the secondary electron detector   (a); accelerating voltage of 1 kV.   Fig. 7. Microstructure and elemental composition of the peripheral region of the surface   oxidised by supersonic air flow for the sample of (HfB2-30vol%SiC)-10vol%Ta4HfC5: in   Journal Pre-proof  27     \\x0c', 'Journal Pre-proof  Fig. 8. Distribution of Hf and Si elements on the oxidised surface of the sample of (HfB2 30vol%SiC)-10vol%Ta4HfC5, peripheral region.   Fig. 9. X-ray pattern of   the oxidised   surface   for   the   sample of   (HfB2-30vol%SiC) 10vol%Ta4HfC5 (top) and X-ray pattern of the Hf6Ta2O17 phase simulated according to the   structural data [86] (bottom).   28       \\x0c', 'Journal Pre-proof  29         \\x0c', 'Table 1. The average surface temperature of samples based on the HfB2-SiC system under   exposure to a supersonic flow of dissociated air, depending on the thermal load1: N - generator anode   power supply, q heat flux   Exposure   parameters   N,   kW   q,   W\\uf0d7cm-  2   Average surface temperature (pyrometer), \\uf0b0С   (HfB2 HfB2 HfB2 30vol%SiC) 30vol%SiC,   30vol%SiC,   HfB2-65vol%SiC,   10vol%Ta4HfC5,   porosity 9%   porosity   porosity 34.5% [4]   porosity 18.5%   [5]   10.8% [6]   (HfB2 30 об.% SiC) 5 об.% Y3Al5O12,   porosity 5.5 %   [84]   1238   1397   30   40   50   60   70   779   1846\\uf0ae1914   (~25 min)2   1515   1590\\uf0ae1623   1433\\uf0ae1487   1555\\uf0ae1848   2620\\uf0ae2560   2300\\uf0ae2554    1640\\uf0ae1740   1712\\uf0ae1754   1408\\uf0ae1405   1490\\uf0ae1461   (25 min)2   (25 min)2   (23 min 30 s)2   1329\\uf0ae1349   1155\\uf0ae1150   1398\\uf0ae1393   363   484   598   691   (25 min 20 s)2   1770\\uf0ae2625\\uf0ae2590   1506\\uf0ae1513   1503-1510   1951\\uf0ae2543   1800\\uf0ae2586   1731\\uf0ae1824   1709\\uf0ae1774   1644\\uf0ae1632   1625\\uf0ae1617   1Pressure in the plasmatron pressure chamber in all experiments was 12-14 hPa;   2The holding time at the maximum anode power supply of the plasmatron of 70 kW.   Journal Pre-proof  30       \\x0c']"
},{
  "_id": 170,
  "PDF": "Oxidation of Porous HfB2–SiC Ultra-High-Temperature Ceramic Materials Rich in Silicon Carbide (65 vol _) by a Supersonic Air Flow.pdf",
  "Text": "['ISSN 0036-0236, Russian Journal of Inorganic Chemistry, 2020, Vol. 65, No. 4, pp. 606-615. © Pleiades Publishing, Ltd., 2020. Russian Text © The Author(s), 2020, published in Zhurnal Neorganicheskoi Khimii, 2020, Vol. 65, No. 4, pp. 564-573.  INORGANIC MATERIALS AND NANOMATERIALS  Oxidation of Porous HfB2-SiC Ultra-High-Temperature Ceramic  Materials Rich in Silicon Carbide (65 vol %) by a Supersonic Air Flow  E. P. Simonenkoa, *, N. P. Simonenkoa, A . N. Gordeevb, A . F. Kolesnikovb, A . S. Lysenkovc, I. A . Nagornova, d, V. G. Sevast’yanova, and N. T. Kuznetsova  aKurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences, Moscow, 119991 Russia bIshlinskii Institute of Problems of Mechanics, Russian Academy of Sciences, Moscow, 119526 Russia cBaikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Moscow, 119991 Russia dMendeleev University of Chemical Technology of Russia, Moscow, 125047 Russia *e-mail: ep_simonenko@mail.ru  Received November 15, 2019; revised November 26, 2019; accepted November 27, 2019  Abstract—Porous HfB2-65 vol % SiC samples (porosity 34.5%) were produced by reactive hot pressing of HfB2-(SiO2-C) composite powder at 1800°C (heating rate 10 deg/min, holding duration 15 min) and 30 MPa. Using a high-temperature induction plasmatron, their resistance to oxidation by a supersonic dissociated air f low was studied (the heat f luxes in the course of the experiment were varied from 363 to 779 W/cm2). The observation of the temperature distribution over the surface of the sample during the experiment showed that a sharp increase in temperature from ~1770-1850 to 2600°C in the samples under investigation occurred at lower heat f luxes and shorter treatment times than that in denser HfB2-30 vol % SiC samples (porosity 9- 11%). This indicated that increasing the density of the HfB2-SiC material and also increasing the silicon carbide content reduced the oxidation resistance. However, the fact that the studied sample withstood 37-min exposure to a high-enthalpy dissociated air f low (including 27 min at a surface temperature of 2560-2620°C) without destruction or complete oxidation makes it possible to assign it to ultra-high-temperature materials.  Keywords: UHTC, thermochemical action, high-enthalpy air f low, induction plasmatron  DOI: 10.1134/S00360236200 40191  INTRODUCTION  Investigations aimed at creating the most eff icient ZrB2(HfB2)-SiC  ultra-high-temperature  ceramic materials (UHTC) remaining functional during aerodynamic heating to ~2000-2500°C are conventionally related to the problem of designing materials for sharp edges of wings and nose parts of high-speed aircraft [1-16]. The great interest in these materials is due to a successful combination of such properties as high values of the melting points, emissivities, and thermal conductivities of the components; relatively high resistance of the materials to oxidation in air f lows (including those containing atomic oxygen); and mechanical properties, fairly good for ceramics. Nonetheless, the strength characteristics, especially the fracture toughness, of UHTC are insuff icient to solve applied problems. Therefore, extensive studies are currently underway to determine the effect exerted on the properties of the produced ceramics by various additives: refractory carbides [17-21], silicides [12, 22], and metal nitrides  [23-25]. Moreover, much  attention has recently been paid to the modif ication of UHTC by metals: chromium [26], nickel [27], vanadium [28], titanium [13, 29-31], aluminum [32], and others. Par ticularly fruitful was the modif ication by various carbon species, f irst of all, carbon nanotubes [33-35] and graphene  [36-39]. The  idea of  reinforcement of ZrB2(HfB2)-SiC  ceramics with  carbon  f ibers  to increase their fracture toughness [40-42] is gradually transforming into the development of technologies for creating ceramic-matrix composites with a boride- carbide matrix [43-47]. The set of useful properties of UHTC materials based on ZrB2(HfB2)-SiC also makes them promising for manufacturing ceramic parts for engines [48, 49]. The optical properties and high thermal conductivity of zirconium and hafnium borides make ceramics based on them good materials for solar absorbers [50- 54]. Owing to the elevated electrical conductivity (also at high temperatures), such materials are quite promising for fabricating ceramic fuel cells [55, 56]. These applications require that the materials should have suff iciently high porosity [50-56]. It is well known that the porosity adversely affects both the mechanical characteristics of a ceramic, and its oxidation resistance. Nonetheless, many attempts were made to increase the thermal-shock resistance and the fracture toughness by forming controllable  606  \\x0c', 'OXIDATION OF POROUS HfB2-SiC ULTRA-HIGH-TEMPERATURE  607  I  2  1  20  HfB2 SiC HfO2 (mon)  40  2θ, deg  60  80  Fig. 1. X-ray powder diffraction patterns of the surface of the HfB2-65 vol % SiC sample (1) before and (2) after exposure to a supersonic dissociated air f low.  porosity [57-64]. It was also shown [60-62] that the above characteristics not always linearly increase with increasing porosity, and the pore size distribution and the  production method  are  important. Besides, decreasing the pore size increases signif icantly the residual strength. For dense ZrB2(HfB2)-SiC ceramics, it is known that the residual strength can be increased by preliminary surface oxidation [63, 64]. Similar studies were carried out for porous materials [65]: ZrB2-SiC samples with various porosities were oxidized in air at 1200°C, and then their residual strength was investigated after testing from 325-725°C in a water bath (25°C). It was found [65] that the thermal shock critical temperature difference Δ Tc of porous samples after oxidation exceeded  that of nonoxidized  samples, which agrees with general concepts. However, there was a reverse dependence of Δ Tc of samples after oxidation on their initial porosity; this was attributed to greater depth of oxidation, which is characteristic of samples with higher porosity. This problem could probably be solved by the laserinduced surface oxidation of porous ZrB2-39 mol % -SiC ceramic composite, which was produced by pressureless sintering at 1900°C for 2.5 h [66]. The laser surface treatment was shown to give rise to a continuous and homogeneous glassy layer ~8 μm thick, which can perform protective functions in oxidation of porous materials. Thus, it can be concluded that the investigation of the oxidation resistance of porous ceramic materials based on zirconium or hafnium diboride doped with silicon carbide is quite an urgent and practically significant problem because it allows one to evaluate their applicability under anticipated operating conditions. Although the thermal behavior of ZrB2(HfB2)-SiC  UHTC in static air was described in many works [3, 26, 67-71], high-speed aerodynamic heating is more correctly modeled using oxygen-acetylene torches, and also electric-arc heaters,  induction plasmatrons, or hypersonic ramjet [2, 5-8, 15, 16, 43, 45, 46, 72-77]. The purpose of this work was to study the effect of a supersonic dissociated air f low on HfB2-SiC ultra-hightemperature ceramic material rich in silicon carbide.  EXPERIMENTAL  HfB2-65  vol %  SiC  ultra-high-temperature ceramic material was produced using a Thermal Technology HP20-3560-20 hot press by reactive sintering of HfB2-(SiO2-C) composite powder [15, 16, 69, 73- 75]. The hot pressing was carried out at a temperature of 1800°C, a holding duration of 15 min, a heating rate of 10 deg/min, and a uniaxial pressure of 30 MPa. The experiment on the exposure of the surface of the porous sample to a supersonic dissociated air f low was conducted using a VGU-4 100-kV high-frequency induction plasmatron. The sample was introduced to a plasma jet at the minimum value 30 kW of the plasmatron anode feed power N. Then, N was stepwise increased from 30 to 60 kW at 10 kW intervals, with the holding duration at each step at N = 30-60 kW being 3 min. Next, the power was increased to N = 67 kW (duration 1.5 min), and at N = 70 kW, the sample was held until the completion of the experiment. The total exposure duration was 37 min. A sonic nozzle with an exit diameter of 30 mm was used. The distance from the nozzle to the sample was 25 mm, the air f low rate was 3.6 g/s, and the pressure in the chamber was 13- 14 hPa. The averaged temperature of the surface of the sample was determined with a Mikron M-770S pyrometer in the spectral ratio pyrometer mode (temperature range 1000-3000°C, measurement spot size ~5 mm). The temperature distribution over the surface of the sample was recorded with a Tandem VS415U thermal imager; in the measurements, the spectral emissivity ε  at a wavelength of 0.9 μm was set to be 0.65 [8, 15, 16, 73-76]. X-ray powder diffraction analysis of the samples before and after exposure to a dissociated air f low was performed with a Bruker D8 ADVANCE X-ray pow radiation, resolution 0.02°). der diffractometer (CuKα Scanning electron microscopy was carried out with a Carl Zeiss NVision 40 focused ion beam scanning electron microscope. The elemental composition of microregions was determined with an Oxford Instruments energy-dispersive X-ray analyzer.  RESULTS AND DISCUSSION  The reactive hot pressing of HfB2-(SiO2-C) composite powder, which was synthesized by the sol-gel method [69, 78], produced a porous ceramic material containing nanocrystalline silicon carbide. The aver RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020  α \\x0c', '608  SIMONENKO et al.  T, °C  2600  N, kW  2400  2200  2000  1800  1600  1400  1200  1000  C  °  0 6 8  0  5  10  15 20 Time, min  25  30  35  pch, hPa 22  20  18  16  14  12  90  80  70  60  50  40  30  Fig. 2. Change with time in the average temperature t of the surface of the porous HfB2-65 vol % SiC sample and in the process parameters: the anode feed power N and the pressure pch in the plasmatron chamber.  age size of crystallites as estimated by the Scherrer formula was 60 ± 2 nm. The density of the samples was 3.93 g/cm3, which corresponded to 65.5% of the theoretical value calculated by the additive method (where the densities of HfB2 and SiC were taken to be 11.2 [79] and 3.2 g/cm3 [80], respectively).  The X-ray powder diffraction analysis showed (Fig. 1, curve 1) that the sample contained no impurity crystalline phases (unreacted SiO2, ZrC, a byproduct of the carbothermal synthesis; or HfO2, a product of the oxidation of hafnium diboride);  there were only phases of hexagonal HfB2 [81] and cubic SiC [82]. The complete conversion of silicon dioxide into carbide directly by hot pressing of HfB2-(SiO2-C) composite powder was also conf irmed by IR diffuse ref lectance spectroscopy data. In particular, after the synthesis, the intense broad absorption band in the wavenumber range 950-1200 cm-1, which represented the stretching vibrations of the Si-O group, vanished, and the absorption band at 700-900 cm-1, which characterized ν (Si-C), emerged. The scanning electron microscopy conf irmed the uniformity of the mutual distribution of grains of HfB2 and synthesized SiC.  To  in situ study the oxidation of the obtained HfB2-65 vol % SiC ceramic material by a supersonic dissociated air f low, a sample was installed into an assembled copper model [15, 16, 73-76]. A mandrel, into which the sample was put, was inserted with a sliding f it into a water-cooled holder. The thermal contact of the end faces of the parts of the model was ensured by the spring action of a tension pin; the con tact surfaces were greased with a thermal paste to improve the heat transfer from the mandrel to the water-cooled holder. To prevent the contact of the sample with the copper mandrel and reduce the heat transfer, the sample was mounted with the end face raised about ~1 mm above the face surface of the model using three SiC whisker wool f ilaments.  The sample was introduced to a dissociated air jet at a plasmatron anode feed power of 30 kW, which was  Time, min  N, kW  q, W/cm2  Table 1. Change with time in the average temperature t of the surface of the HfB2-65 vol % SiC sample in the vicinity of  the critical point (as determined by spectral ratio pyrometer), in the process parameters (the anode feed power N and the pressure in the plasmatron chamber, pch = 13 ± 1 hPa1), and also in the corresponding heat f luxes1 q T, °C 1408 →  1405 1515 1640 1740 1770 →  2625 2590 2615 →  2600 2620 →  2605 2585 2570 2570 2560 2560  363 484 598 598 691 691 760 779 779 779 779 779 779  12 →  13.5 →  20 25 30 35 37  30 40 50 50 60 60 67 70 70 70 70 70 70  9.5 10 13.5 15  0 →  3 →   3 6  6 8  9 →   1The heat f luxes to the water-cooled copper calorimeter were determined in individual experiments described before [83].  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020  \\x0c', 'OXIDATION OF POROUS HfB2-SiC ULTRA-HIGH-TEMPERATURE  609  9 min 0 s  1600 °С  1300  12 min 59 s  691 W/cm2  1000  760 W/cm2  2800  °С  2000  1400  14 min  779 W/cm2  2800  °С  2000  1400  6 min 24 s  1600 °С  1300  598 W/cm2  1000  3 min 20 s  484 W/cm2  0 min 36 s  363 W/cm2  1500 °С  1200  900  1400 °С  1150  900  2500  2000  C  °  ,  T  1500  1000  -10  0 Radius, mm  10  20 min  779 W/cm2  30 min  779 W/cm2  37 min 1 s  779 W/cm2  37 min 5 s  779 W/cm2  2800  °С  2000  1400  2800  °С  2000  1400  2800  °С  2000  1400  2200  °С  1700  1200  Fig. 3. Temperature distribution over the surface of the sample at various stages of exposure to a supersonic dissociated air f low and also along the diameter (black straight line).  then stepwise increased to 70 kW. Table 1 and Fig. 2 present the interaction parameters and the average temperatures of the surface of the samples as determined by the spectral ratio pyrometer.  It is seen from Table 1 and Fig. 2 that, at a plasmatron anode feed power of as low as 40 kW (3-6 min of experiment, q = 484 W/cm2), there was a weak tendency for the surface temperature to rise with time. This tendency is much stronger at the second step, when the average surface temperature increased in 3 min by 130°C from 1640 to 1770°C. At N = 60 kW (q = 691 W/cm2), the average surface temperature in under a minute sharply rose to 2625°C. A further increase in the plasmatron anode feed power to 67 and 70 kW led only to an insignif icant increase in the temperature of the surface of the sample and its subsequent decrease to a stable value of 2570-2600°C.  For more detailed investigation of the processes that occurred on the surface of the material and in its  surface layers and led to the above change in the average temperature, the surface temperature distributions at various stages of the experiment were recorded with a thermal imager (Figs. 3, 4). It was shown [15, 16, 74] that, on exposure to a supersonic f low, unlike a subsonic one, even at relatively low heat f luxes (363-598 W/cm2), there was  no  uniform  temperature  distribution throughout the surface of the sample being oxidized. With increasing plasmatron anode feed power, a hot spot formed at the center of the sample, rather than at its periphery, as in the case of heating by a subsonic air f low.  It was from the overheated central region that the high-temperature (>2000°C) spot grew (Figs. 4, 5). Whereas at the f irst stages (N = 30-50 kW) the temperature difference between points at the center and periphery of the sample was ~50°C, after a sharp and very rapid (in 40 s) increase in the average temperature, this difference was as large as 160-260°C (Fig. 4).  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020    \\x0c', '610  SIMONENKO et al.  2800  2600  2400  2200  2000  1800  1600  1400  1200  1000  C  °  ,  T  2690°C  °  5 2 2  °  0 1 2  2540°C  °  0 6 2  Point 1  Point 2  1  2  °  0 6 1  2400 °С  2000  1600  1200  0  5  10  15  20  25  30  35  Time, min  Fig. 4. Change in the temperatures of the surface of the sample in the course of the exposure at the center (point 1, red curve) and periphery (point 2, blue curve) of the sample.  Note that the rate of heating of the entire surface of the sample from 1750-1800 to 2600-2500°C is very high even in comparison with that observed under the action of a supersonic f low on the HfB2-SiC samples that were obtained by a similar method, but contained less silicon carbide (30 vol %), and had lower porosity. Figure 6 presents the superimposition of the initial portions of the curves of the change in the average surface  temperature  (according  to  the  spectral  ratio pyrometer data) for the HfB2-65 vol % SiC sample (porosity 34.5%) studied in this work and the HfB2- 30 vol % SiC samples with somewhat different porosities (10.8% [15] and 9.0% [16]). The variations of the anode feed power and the pressure in the plasmatron chamber in all the three cases coincided within the experimental error. One can see that the sharp temperature increase from 1750-1800 to 2600-2700°C1 for the more porous sample, which was also richer in SiC, occurs at a lower heat f lux than that for the denser HfB2-30 vol % SiC samples (q = 691 W/cm2 as against 779 W/cm2 [15, 16]). There were also differences in behavior between the samples of identical compositions (30 vol % SiC): the heating duration required for “surface temperature jump” [84] was longer for the denser HfB2-30 vol % SiC sample (Fig. 6).  Furthermore, the rates of heating of the surface from 1770-1820 to 2600-2700°C also differed signif i 1 Because of the evaporation of the silicate glass layer and the appearance of porous, highly catalytically active, and low thermally conductive HfO2 on the surface [8, 15, 16, 73-76, 83].  cantly: for the densest HfB2-30 vol % SiC sample (porosity 9.0% [16]), this rate was ~9 deg/s; for the sample of the same composition with a porosity of 10.8% [15]; it was ~11 deg/s; and for the for the HfB2- 65 vol % SiC sample studied in this work (porosity 34.5%), the rate is almost twice as high: 21 deg/s. Thus, along with the composition of an ultra-hightemperature ceramic material, the conditions and rate of evaporation of silicate glass from the surface were signif icantly affected by the porosity. Although the porous HfB2-SiC sample rich in silicon carbide (65 vol %) was oxidized more intensely, it withstood 40-min exposure to a supersonic dissociated air f low without destruction. After switching-off of heating, no thermal-shock cracking of the sample while fast cooling was observed. The total mass loss after 37 min of exposure was 8.6%, which considerably exceeded Δ m for the previously studied [15, 16] denser HfB2-30 vol % SiC samples (2.6-3%). The increase in the thickness of the cylindrical HfB2-65 vol % SiC sample after the oxidation was ~3%. The X-ray powder diffraction analysis (Fig. 1, curve 2) demonstrated that the main crystalline product of the oxidation is monoclinic hafnium oxide. The scanning electron microscopy showed (Fig. 7) that the oxidation gave rise to a typical porous ceramic structure, which, in this case, was more voluminous. The microstructures of the central and peripheral regions were similar; however, on the surface at the edges of the sample, there were traces of silicate glass.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020    \\x0c', '2400 °С  1800  1200  2400 °С  1800  1200  2400 °С  1800  1200  1800 °С 1400  1100  1800 °С 1400  1000  9 min 22 s  9 min 20 s  9 min 18 s  9 min 14 s  9 min 11 s  2600  2400  2200  C  °  ,  T  2000  1800  1600  1400  9 min 30 s  9 min 32 s  9 min 34 s  2500 °С  1900  1300  2600 °С 2000  1300  2600 °С 2000  1400  2600 °С 2000  1400  2600 °С 2000  1400  OXIDATION OF POROUS HfB2-SiC ULTRA-HIGH-TEMPERATURE  611  2400 °С  1800  1200  9 min 24 s  2400 °С  1800  1200  2500 °С  1900  1300  9 min 28 s  9 min 26 s  -10  0 Radius, mm  10  9 min 36 s  1600 °С 1300  1000  9 min 0 s  2600 °С 2000  1400  9 min 40 s  9 min 38 s  Fig. 5. Temperature distribution over the surface of the sample at the time of the so-called “temperature jump” (9-10 min of exposure) and also along the diameter (black straight line).  The energy-dispersive analysis determined that, on the average, n(Hf) : n(Si) = 1 : 14.5; i.e., on the surface, there was primarily hafnium oxide, as it was observed for most of the samples that were heated for a long time to 2500-2700°C [8, 15, 16, 74-76, 83].  CONCLUSIONS  The porous HfB2-65 vol % SiC ceramic material produced by reactive hot pressing was oxidized by a supersonic dissociated air f low using a high-temperature induction plasmatron.  The study of the oxidation of the ceramic determined that the main stages of the oxidation are the same as those described earlier for ultra-high-temperature boride-carbide ceramics:  (1) the surface oxidation of HfB2 and SiC;  (2) the formation of a protective molten silicate layer on the surface;  (3) the further oxidation of the HfB2-SiC ceramic in the bulk of the sample by oxygen that has diffused through the silicate glass layer (silicon carbide is oxidized by an active mechanism to form gaseous SiO);  (4) under certain experimental conditions (a combination of certain values of the heat f lux, the pressure in the plasmatron chamber, and the exposure duration), the surface temperature reaches ~1750-1850°C;  (5) the evaporation of silicate glass from the surface of the melt, after which porous, highly catalytically active, and low thermally conductive hafnium dioxide appears on the surface; at this stage, there is a “temperature jump” to ~2500-2700°C;  (6) the stabilization of the surface temperature and certain (relatively low) contents of silicon and boron  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020    \\x0c', '612  SIMONENKO et al.  65 vol % SiC, porosity 34.5% this work  30 vol % SiC, porosity  9.0% [16]   30 vol % SiC, porosity 10.8% [15]   C  °  ,  T  2500  2000  1500  1000  0  5  10 Time, min  15  20  Fig. 6. Superimposition of the initial portions of the curves of the change in the average surface temperature on exposure to a supersonic dissociated air f low (at similar heat f luxes) for the HfB2-65 vol % SiC sample and the HfB2-30 vol % SiC samples with somewhat different porosities [15, 16].  Central region  Peripheral region  100 μm  100 μm  500 nm  4 μm  20 μm  20 μm  Fig. 7. Microstructure of the oxidized surface of the HfB2-65 vol % SiC sample in the central and peripheral regions.  oxides in the in the gas phase in the boundary layer over the sample. This suggests that the borosilicate glass layer located deeper in the material continues to perform its protective function and prevents the complete oxidation of the sample despite its high porosity (34.5%).  Nonetheless, it was noted that the sharp increase in the temperature from 1770-1800°C to ~2600°C for the porous HfB2-65 vol % SiC sample occurs at lower heat f lux and shorter total duration of exposure to a supersonic dissociated air f low than that for the similarly obtained HfB2-30 vol % SiC  samples with  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 4    2020    \\x0c', 'OXIDATION OF POROUS HfB2-SiC ULTRA-HIGH-TEMPERATURE  613  a porosity of 9-11% [15, 16]. Besides, in this work, a temperature on the order of 2600°C was established throughout the surface of the sample for a shorter time (in 40 s). Summarizing the obtained results, one can conclude that porous HfB2-SiC materials rich in silicon carbide (65 vol %) are much less resistant to oxidation. 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},{
  "_id": 171,
  "PDF": "Oxidation of spark plasma sintered ZrC-Mo and ZrC-TiC composites.pdf",
  "Text": "[\"International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  Contents lists available at ScienceDirect  International Journal of Refractory Metals & Hard Materials  jou rna l homepage : www . e lsev ie r .com / loca te / I JRMHM  Oxidation of spark plasma sintered ZrC-Mo and ZrC-TiC composites  Der-Liang Yunga,⁎, Birgit Maatenb, Maksim Antonova,  Irina Hussainovaa,c,d  a Department of Materials Engineering, Tallinn University of Technology, Ehitajate tee 5, 12616 Tallinn, Estonia b Department of Thermal Engineering, Tallinn University of Technology, Ehitajate tee 5, 12616 Tallinn, Estonia c ITMO University, Kronverksky 49, St. Petersburg 197101, Russian Federation d Department of Mechanical Science and Engineering, University of Illinois at Urbana-Champaign, 1206 West Green Street, Urbana,  IL 61801, USA  MARK  A R T I C L E  I N F O  A B S T R A C T  Keywords: Spark plasma sintering Oxidation Ceramic composites ZrC-Mo ZrC-TiC  Two ZrC-based composites, ZrC-20 wt% Mo cermet and ZrC-20 wt% TiC solid solution mixed carbide, were prepared by spark plasma sintering. A sample of pure ZrC was also included in the test to be a benchmark reference. The oxidation performance of the composites was studied after exposing the samples to temperatures between 600-1200 °C in air. Thermogravimetric analysis was made to determine the weight change during oxidation, along with using XRD analysis to profile the chemical composition of the oxide layer at each temperature. Both pure ZrC and ZrC-Mo cermet suffer catastrophic oxidation at 1200 °C, either undergoing spalling or pesting, reducing the material to flakes or powder. ZrC-TiC, however, was able to resist severe oxidation damage up to 1200 °C. The oxide layer was determined to contain mixed oxide species (Zr,Ti)O2 attributed to the mixed carbide, solid solution nature of the ZrC-TiC composite. This mixed oxide species was able to exert a more protective effect on the overall matrix beneath the oxide layer, stalling deeper oxidation into the microstructure.  1.  Introduction  The class of ultra-high temperature ceramics (UHTCs) such as zirconium or titanium carbides (ZrC or TiC) are widely investigated as materials for high-temperature structural applications due to their high melting points and retained mechanical strength even at up to 3000 °C under protective environments [1]. ZrC is regarded as an important candidate material as it has received attention as either an additive or a base material for composite alloys. ZrC possesses a high melting pointing (~ 3420 °C), high hardness (~ 25.5 GPa), low elec(7.8 × 10− 7 Ω·cm), trical resistivity and high modulus of elasticity (~ 400 GPa) [2]. A common application for ZrC alloys exists in the nuclear industry as cladding shields for radioactive materials since ZrC's ionic bonding is highly resistant against irradiation damage [3,4]. However, it has been well documented that pure ZrC has poor hightemperature chemical stability against an oxidising atmosphere [5,6], thus limiting its potential application as a UHTC. ZrC oxidation begins as low as 600 °C depending on oxygen partial pressure, where a c-ZrO2 compact phase is formed preventing further oxidation of carbon in the ZrC [7,8]. At higher temperatures beyond 800 °C, with higher a partial pressure of oxygen, ZrC quickly oxidises into either monoclinic and/or tetragonal ZrO2 oxide scales consisting of a porous and cracked outer layer [9]. The resulting oxides offer virtually no protection against  further oxidation, confirmed by extrapolated linear oxidation kinetics [10]. Various avenues of research exist to alloy ZrC with other compounds and elements to enhance the oxidation resistance of ZrC. One of the compounds alloyed to ZrC to enhance oxidation resistance is silicon carbide (SiC); the resulting silicon oxide (SiO2) is supposed to form a more protective oxide layer. However, the available research seems to portray conflicting results concerning ZrC-SiC composites. Zhao et al. [11] tested ZrC-30 vol% SiC in air between 800 and 1500 °C, but found that the resulting SiO2 only offered protection at a higher temperature range, above 1100 °C. In contrast, Ma et al. [12] experiments with ZrC-20 vol% SiC claimed that oxidation resistance, in accordance with mass change, was better below 1000 °C. The kinetics of ZrC alloy oxidation seems to be dependent on the concentration of the alloying compound, SiC. A later study [5] with ZrC-SiC found that higher concentrations (~ 30 wt% SiC) exhibited better ZrC oxidation protection at lower temperatures (~ 1500 °C) compared to lower concentration (~ 10 wt% SiC) performance, offering better protection at higher temperatures (~ 1800 °C). The variables at play during oxidation are complex and numerous since most experiments are done in static air conditions where variables, such as partial pressure of oxygen that can influence oxidation kinetics, are not taken in account. Lower oxygen partial pressure improved the oxidation resistance of ZrB2-SiC ceramics for example [13].  ⁎ Corresponding author at: Tallinn University of Technology, Ehitajate tee 5, Tallinn 19086, Estonia. E-mail address: derliangyung@gmail.com (D.-L. Yung).  http://dx.doi.org/10.1016/j.ijrmhm.2017.03.019 Received 28 June 2016; Received in revised form 22 March 2017; Accepted 27 March 2017  Available online 06 April 2017 0263-4368/ © 2017 Elsevier Ltd. All rights reserved.  \\x0c\", \"D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  The focus of this paper involves two composites: zirconium carbide molybdenum (ZrC-Mo) and zirconium carbide-titanium carbide (ZrCTiC). Research had been done on producing ZrC cermets with the goal to enhance the fracture toughness of ZrC. Landwehr et al. [3,14] made extensive studies on ZrC-Mo cermet production under various sintering conditions. Their findings concluded that successful incorporation of Mo into ZrC decreased the cermet's thermal expansion coefficient and improved thermal shock resistance. The authors of this paper continued their research employing spark plasma sintering (SPS) technology for ZrC-Mo synthesis [15]. We found that the mechanical properties of ZrCMo cermet were influenced by sintering temperature rather than compaction pressure. Similarly using SPS technology, ZrC-TiC has also drawn research attention towards the composite's mechanical properties and manipulating the miscibility gap for solid-solution formation [16]. We found that increasing compaction pressure during SPS up to 100 MPa at 1600 °C produced ZrC-TiC with distinct carbide phases, yielding better fracture toughness values than the same composites produced at higher temperatures [17]. Given the promising mechanical and thermal properties of these two composites, there is an evident lack of research done on their oxidation characteristics at elevated temperatures. Research has been done on ZrC oxidation as well as the oxidation of other elemental composites, including Mo [18] and TiC [19] base. However, the oxidation behavior results of ZrC-20 wt% Mo (~ 14 vol% Mo) and ZrC-20 wt% TiC (~ 25 vol% TiC) composites in an air furnace, at high temperatures up to 1200 °C, have not been done. This article looks to examine the microstructural evolution of oxide formation, analysed on the surface and cross sections of the oxidised samples. Scanning electron microscope (SEM), thermogravimetric analysis (TGA), and X-ray diffraction (XRD) results are also included in the analysis with the TGA giving information about weight change during oxidation.  2. Materials and methods  (~ 3.6 μm, Pacific Particulate Commercially available ZrC powder Materials, 99% purity ZrC-7643) underwent ball milling with either 20 wt% Mo (1.0-3.0 μm, Pacific Particulate Materials, 99% purity Mo(2-3 μm, Pacific Particulate Materials, 7164) or 20 wt% TiC powder 99% purity TiC-2088). The compositions were milled for 48 h in a tungsten carbide mill using TiC-NiMo balls in a 8:1 wt% ball to powder ratio along with ethanol as a milling liquid. The powders were dried for 24 h at 50 °C, then sieved to 200 μm. Spark plasma sintering machine (FCT Systeme GmbH) was used to synthesis all samples under vacuum with parameters outlined in Table 1 according to outlines specified in previous research [15,17]. ZrC needed to be SPS processed under nitrogen gas (approx. 60 mbars), as the temperature exceeded 2000 °C, to maintain the integrity of the graphite mould. One deviation from the previous research was to adopt control, slow cooling for the samples after sintering from the final temperature to 300 °C at a rate of ~ 22 °C ∗ min− 1 to prevent thermal shock cracks during post processing. The samples were mechanically ground and polished with diamond abrasives down to a 1 μm finish. Then the disc samples, measuring 5 mm thick by 20 mm diameter, were cleaned via ultrasonication acetone bath and dried, before being weighed using an electronic balance with accuracy 0.0001 g. Oxidation tests were carried out in a Nabertherm static air furnace  Table 1 SPS parameters for all samples according to [15,17].  Sample  Ramp [°C ∗ min− 1]  Temp [°C]  Dwell [s]  Pressure [MPa]  ZrC ZrC-Mo ZrC-TiC  100 100/300 100  2200 1600 1600  300 300 600  50 50 100  in laboratory air between 600 °C and 1200 °C, at intervals of 200 °C. The experiment was setup so that separate samples were exposed at each temperature point only once before cooling via ambient atmosphere. The furnace ramp rate was calibrated to be 12 °C ∗ min− 1 for a 1 h dwell time at each temperature point. The microstructures of the oxidised surface and cross-sections were examined under scanning electron microscopy (FE-SEM Hitachi S-4700, Japan) equipped with energy dispersive spectroscopy (EDS-TM-1000) analyser. Surface analysis via SEM was done as-is, while cross-section microstructural 1 μm finish with analysis was done after polishing to a diamond abrasive. Chemical composition was analysed with X-ray diffraction analysis (XRD, Philips PW3830 X-ray Generator, 4 kW, Cu-Anode) using CuKα radiation. The accelerating voltage was 40 kV with a filament current of 30 mA, a scan step size of 0.02°, and a count time of 0.4 s at each step. Each composite underwent thermogravimetric analysis using a NETZSCH STA 449 F3 Jupiter® thermal analyser with differential scanning calorimetry (DSC). The system contains a high heating rate furnace, which is capable of operating between 30 and 1250 °C. The system is air tight, allowing measurements to be conducted under precisely defined atmospheres. The software, NETZSCH Proteus 6.1.0, allows the computation of the rate of mass change, mass change steps, onset and peak temperature, peak area integration etc. The samples were analysed in Al2O3 crucibles without lids. The heating rate was in all cases 20 °C ∗ min− 1. The atmosphere was 20% O2 and 80% N2 to simulate the effect of ordinary air.  3. Results and discussions  3.1. Weight and thermogravimetric analysis  The mechanical properties, densities, microstructure, and XRD analysis of the composites ZrC-20 wt% Mo and ZrC-20 wt% TiC at room temperature are described elsewhere [15,17]. Fig. 1 shows the weight gain and loss at the oxidation temperatures for ZrC-Mo, ZrC-TiC, and ZrC samples at 600 °C, 800 °C, 1000 °C, and 1200 °C respectively. All samples show negligible weight change at 600 °C suggesting oxidation evolution is kept at a minimum. Both ZrC-Mo and ZrC-TiC show similar oxidation weight gain at 800 °C, while the reference ZrC exhibits significant oxidation weight gain above 600 °C. This result for ZrC oxidation has been well documented as oxygen penetrates deeper into the microstructure, oxide formation is propagated at this temperature range [7-9]. From Fig. 1, it can be concluded that ZrC-Mo and ZrCTiC perform better than pure ZrC at resisting oxidation up to 800 °C, or at least halts the deeper penetration of oxide formation into the bulk structure. However, as the temperature rises above 800 °C, this is where  Fig. 1. Weight gain per surface area of ZrC, ZrC-Mo, and ZrC-TiC composites plotted for  1 h dwell at each temperature.  245  \\x0c\", 'D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  Fig. 1 and Fig. 2A is likely due to the sample size and the timing of the experiments. A larger sample in the static air furnace means less surface area to volume ratio, thus oxidation takes longer to complete. This is compared to the tiny pieces used in the TGA machine, which oxidised completely and hence the prolonged weight plateau beyond 800 °C. The TGA graphs largely support the generalised trend established in Fig. 1. ZrC-Mo is able to hold off MoO3 volatilisation until the temperature ramps above 800 °C. Afterwards there is catastrophic weight loss of the sample seen in Fig. 2B during temperature ramping to 1000 °C. MoO3 volatilising into gas causes the rapid weight loss before stabilising again at 1200 °C. The kinetics indicate a paralinear oxidation trend for ZrC-Mo, but in a negative matter where weight loss is most rapid during temperature rising. Fig. 2C reveals ZrC-TiC to have more aggressive oxidation taking place as the temperature proceeds above the 1000 °C threshold. Overall the TGA graphs show ZrC-TiC composite to follow similarly along the paralinear oxidation law of pure TiC, where the parabolic character of oxidation dominates in the part of the process [23]. The TGA plots can be visibly quantified in the following images shown in Fig. 3, where photographs were taken after each sample had been tested up to the indicated temperature point then cooled back to the room temperature. To prepare the samples shown in Fig. 3 samples underwent grinding and polishing to yield the same metallic shine finish before oxidation experiments. At room temperature, all samples appear similar. As reflected in the TGA graphs all samples show little oxidation effect at 600 °C. ZrC would be the only exception, as at 600 °C, there is a hint of white oxides appearing on the surface, view macroscopically. All three samples exhibit white oxide formation at 800 °C on their respective surfaces with ZrC sample appearing to have the most oxide thickening especially around the edges of the disc by the occurrence of volume expansion. Cracks were seen with ZrC-Mo and ZrC-TiC at 800 °C were due to thermal stresses incurred during the synthesis and SPS of the sample, aggravated during the grinding and polishing. In subsequent samples, slow cooling after SPS eliminated such breakages in the samples. Of course, cracks in the sample can certainly affect the surface area-volume ratio during oxidation testing, but since the weight data between air furnace and TGA correlate one another, it is then safe to assume this incident was negligible.  3.2. ZrC-Mo composite  At 1000 °C, the three samples start to differentiate in appearance and chemical reaction. The obvious note is the surfaces of ZrC-Mo and ZrC-TiC become a yellowish tinge. The presence of MoO3 has been known to give a light, yellow fume indicative of it melting point, 790 °C [22] through the following reaction:  2Mo + 3O → 2MoO  2  3  (1)  However, when dealing with ZrC-Mo composites, this composites presents a complex chemical reaction since Mo exist as a solid solution with ZrC and also carburises the ZrC carbon forming MoC or Mo2C in the reaction [15,24]:  ZrC + Mo → MoC + (Zr, Mo) C  (2)  As shown in XRD patterns, Fig. 4, ZrC-Mo is mainly composed of (Zr,Mo)C solid solution and MoC in the native microstructure. Therefore, when oxidation takes place for ZrC-Mo, the reactions are as follows:  ZrC + 2O → ZrO + CO  2  2  2 (g)  ZrC + O → ZrO + C  2  2  (s)  2MoC + 5O → 2MoO + 2CO  2  3  2 (g)  (3)  (4)  (5)  The entire surface of ZrC-Mo becomes a mixture of MoO3 and ZrO2; ZrC oxidises to form both tetragonal [ZrO2(t)] and monoclinic [ZrO2(m)] zirconia. It is surprising that MoO3 does not fully volatised  246  Fig. 2. TGA plots percent weight change between 800 and 1200 °C under simulated air  conditions: A) ZrC, B) ZrC-Mo, C) ZrC-TiC.  the oxidation curve for ZrC-Mo exhibits the oxidative nature of Mo. It is well-known that molybdenum and molybdenum alloys are susceptible to losses from the highly volatile molybdenum trioxide (MoO3) species [18,20,21]. Studies have documented steady material weight lost due to MoO3 volatilisation beginning at ~ 800 °C [18,22]; the ZrC-Mo composition is able to resist weight loss at 800 °C. The reasons for this are discussed in the next section. Between 800 °C and 1000 °C, both ZrC and ZrC-TiC plateau in terms of their oxidation weight changes, but ZrC-Mo continuously loses weight. It is important to keep in mind that the Fig. 1 data are drawn from single weight measurements after each oxidation experiment, at each temperature point. A more detailed analysis is drawn from TGA graphs presented in Fig. 2, where the weight is constantly monitored within a controlled environment. Fig. 2A supports the oxidation kinetics of ZrC, where initially the ZrO2 with carbon forms a compact, pore-free microstructure, which acts as an oxygen barrier to for further oxidation between room temperature and 700 °C [8]. Oxidation for ZrC starts from approximately ~ 600 °C and starts to plateaus at 800 °C and beyond, most of the carbon form will be oxidised into gaseous species rendering the protective layer useless [11]. The oxide scales become non-adherent and brittle. The discrepancy between ZrC oxidation in  \\x0c', 'D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  Fig. 3. Photograph after oxidation at different temperatures for 1 h then subsequently cooled to room temperature: A) ZrC, B) ZrC-Mo, and C) ZrC-TiC. The 600 °C and 800 °C images for  all samples show a zoomed in portion of the sample to detail the subtle fretting at the edges. (For interpretation of the references to colour in this figure, the reader is referred to the web  version of  this article.)  pore-pore microstructure that acts as a diffusion barrier against oxygen transport [8] so the local oxygen potential becomes too low to further oxidise the Mo and ZrC. This would explain why MoO3 volatilisation is inhibited until about 900 °C according to the Fig. 2B. The readied formation ZrO2 at lower temperatures of ~ 600 °C effectively acts like a protective shield, shielding the more volatile Mo phases from excessive oxidation and subsequent volatilisation. Of course, ultimately, as the temperature rises towards 1000 °C, there would seem to be a systematic break down of the ZrO2 shield as MoO3 formation overtakes the material surface. The thermal mismatch between ZrO2 and MoO3 also causes further cracking in the microstructure leading to an oxidation cascade effect. The oxidation temperature up to 1200 °C ultimately causes both ZrC and ZrC-Mo to suffer catastrophic structural damage in a phenomena known as pesting: disintegration the material into powder [26]. MoO3 vaporises from the external and cracked surfaces initiate cracks to become porous. Oxygen can then penetrate through the sample into cracks causing volume expansion, furthering oxidation until the material is reduced to powder as seen Fig. 2B, 1200 °C. Given the size of the sample and time exposure at high temperatures, the oxidation penetration did not extend completely through the sample leaving some native ZrC-Mo in the macrostructure. MoO3 is not detectable by XRD above 1000 °C suggesting all the surface traces were volatised leaving behind only ZrO2. Fig. 5 shows sequential SEM images of the oxidation process of ZrCMo between 600 °C and 1200 °C after 1 h at each temperature. The surface of Fig. 5A shows at 600 °C, there is visible depressions of the darker MoC phase and etching around the (Zr,Mo)C grains. The cross section at 800 °C shows a clear separation line between zirconia-rich, darker layer with white spots of MoO3 contrasted to the lighter, unaffected ZrC-Mo layer, Fig. 5B. At 1000 °C, Fig. 5C shows a visible interlayer between the depleted outer ZrO2 and MoO3 layer (Fig. 5D),  Fig. 4. XRD patterns of ZrC-Mo surface at room temperature (RT), and then after 800 °C,  1000 °C, and 1200 °C. Both tetragonal [ZrO2(t)] and monoclinic [ZrO2(m)] zirconia were detected during oxidation.  as would be expected in a typical Mo alloy such as titanium, zirconium molybdenum (TZM) where most oxides would volatilise by 800 °C, losing up to 33% mass [25]. This is not seem with ZrC-Mo at 800 °C as there is no evidence yet of a molten layer of MoO3 indicated by a yellow tinge in Fig. 3B. Notice that the diffraction intensity of MoO3 peaks are subtle at 800 °C, but increase intensity at 1000 °C before disappearing completely at 1200 °C. It is possible the ZrO2 phase forms a compact,  247  \\x0c', 'D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  Fig. 5. SEM images,  including EDS scan, of ZrC-Mo after different oxidation temperatures in static air.  which has already become porous from the MoO3 volatilisation, and unaffected native layer. EDS scan supports the idea that surface layer is composed primarily of ZrO2 and MoO3 oxides, but also hints of Ni, Ti, and W most likely due from the ball milling mediums used to refine the powder. This interlayer between the oxide and unaffected layer (Fig. 5E) is believe to be a part of the mechanism of MoO3 oxidation where at temperatures between 800 and 1100 °C, very thin coatings of MoO2 exist, which then undergoes oxygen absorption via the reaction [25,27]:  O + MoO → MoO  2  3  (6)  The existence of the MoO2 interlayer at 1000 °C would help explain the TGA data presented in Fig. 2B, where there is oxidation weight loss due to MoO3 volatilisation occurs during the ramp rate to a higher temperature but then reverts to paralinear kinetic during the dwell step. It is possible to suggest the stoichiometry of the Mo oxide changes as the scales grew, thus changing the rate of volatilisation process. A thicker outer layer of remaining ZrO2 may have been enough to protect the inner MoO2 layer to slow down its oxidation rate. The final SEM image in Fig. 5F, in conjunction with XRD of Fig. 4, show only particles of ZrO2 remaining on the surface as no Mo oxide is detected.  248  3.3. ZrC-TiC composite  At the macrosurface, unlike what is seen with pure ZrC and ZrC-Mo cermet, ZrC-TiC is abled to resist severe oxidation damage most effectively up to the highest experimented temperature. As Fig. 2C shows at 1000 °C, while ZrC and ZrC-Mo suffer spalling, ZrC-TiC maintains its integrity. From the XRD patterns, Fig. 6, ZrC-TiC is sintered as a composition of (Zr60Ti40)C since the higher pressure of SPS forces TiC dissolution into ZrC base on its miscibility gap producing a mix carbide solid solution [16,28,29]. TiC is not entirely dissolved, however, as some free TiC is still detected in XRD patterns at room temperature. The oxidation of carbides of group IV transition metals occurs through formation of an oxy-carbide of the metal plus carbon [6,8,30], which is subsequently oxidised to CO and CO2, eventually yielding metal oxides according to the reaction (7):  MeC + 2O → MeC O + C  2  x  y  1− x  + O  4− y  → MeO + CO  2  2  (7)  At 800 °C, the first signs of oxidation occur mainly with ZrO2 (from reaction (3)) and TiO2 formation reaction in air. However, taking into account reaction (7), the following reactions are reasonable expectations for ZrC-TiC:  TiO + 2O → TiO + CO  2  2  2  2 (g)  (8)  \\x0c', 'D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  (Zr, Ti) C + 2O → (Zr, Ti) C O + C  2  x  y  1− x  + CO  2 (g)  + O  4− y  → (Zr  , Ti  0.85  ) O  2  0.15  (9)  Much of the ZrC-TiC oxidation reaction would follow the route of reaction (9) with the mixed carbide solid solution oxides, which is seen in XRD scans starting from 800 °C. Nevertheless, there is strong indication for reaction (8), based on the assumption that the TiC is fully stoichiometric [31], does occur with strong diffraction peaks at 27° and 36° identified as rutile TiO2 [32] evident in Fig. 6, 800 °C-1200 °C. Rutile does possess a slight yellow tinge in its native form, but the most likely explanation for the yellowish hue during ZrCTiC high temperature oxidation comes from the yellow pigmentation of monoclinic zirconia as typically happens between 800 and 1300 °C [33] and supported by the peak intensity of monoclinic ZrO2 increases between 1000 °C and 1200 °C in the XRD analyses and in the benchmark ZrC in Fig. 3A, 1000 °C. SEM image details the step-wise evolution of oxide formation for the ZrC-TiC composite, where the initial oxide layer is barely visible forming at 600 °C, Fig. 7A. This thin layer measures approximately 5-6 μm, visible as a darker region in the image, separated by a dotted red line. Oxidation becomes more prevalent as temperatures reach the oxide layer has already extended 200 μm 800 °C since after 1 h; composing of a ZrO2 and TiO2 rich layer, Fig. 7B. This is also the point  Fig. 6. XRD patterns of ZrC-TiC surface at room temperature (RT), and then after 800 °C,  1000 °C, and 1200 °C. Both tetragonal [ZrO2(t)] and monoclinic [ZrO2(m)] zirconia were detected during oxidation.  Fig. 7. SEM images,  including EDS scan, of ZrC-TiC at different oxidation temperatures in static air.  249  \\x0c', 'D.-L. Yung et al.  International Journal of Refractory Metals & Hard Materials 66 (2017) 244-251  where the first evidence of the (Zr,Ti)O2 mixed oxide species is detected by XRD, Fig. 6. At 1000 °C, both images in Fig. 3B and Fig. 7C show loss of adhesion between the oxide and matrix layers. It is unclear why this happens, yet the overall structure of ZrC-TiC shows better integrity than ZrC-Mo or ZrC at 1000 °C and 1200 °C. One reason for the oxide and depleted matrix layer loose adhesion comes simply from the formation of ZrO2 and TiO2 at the intersection of the mixed oxide layer at the surface and unaffected matrix. Similarly to the formation of zircon in a - zirconium diboride silicon carbide composite oxidation, where a tough outer layer of zircon forms above a more porous layer of individual oxides [34], a porous layer of oxide forms between the matrix, which is protected by a dense layer of mixed oxide (Zr,Ti)O2. Conversely, the mixed oxide species, (Zr,Ti)O2, was determined to be tetragonal, hence a mismatch between the cubic (Zr,Ti)C matrix could also explain why the layers adheres so loosely. Another reason maybe the mismatch in coefficient of thermal expansion (CTE) between the oxidised scale and un-oxidised matrix together act as a bilayer structure. Assuming the average linear coefficient of thermal expansion of the oxide scale is greater than that of the carbide-based specimen, then the thermal mismatch stress upon cool down would be tensile at the exterior surface and compressive within the carbide specimen [35]. The cracks would originate at the exterior surface and terminated at the oxide-carbide interlayer as seen in Fig. 7C. The cracks at the surface shown in Fig. 7D and between the bilayer in Fig. 7E show further penetration into the matrix, which in theory should accelerate internal oxidation for the ZrC-TiC sample. However, in our results, this is not the case. In fact, from Fig. 3C 1200 °C, there is clear indication of an oxidation gradient developing between the outer surface (yellow hue) into the depleted zone (lighter yellow hue) and into the matrix (grey). The final image of the oxidised ZrC-TiC surface, Fig. 7F, shows two contrasting grains, white being undisguisable ZrO2 and TiO2, and larger, darker grains to be the (Zr,Ti)O2. The data on (Zr,Ti)O2 is not well known, therefore the exact kinetic nor mechanism that allows ZrCTiC to resist oxidation damage better than ZrC-Mo or pure ZrC has not fully been determined. Nevertheless, other studies involving mixed oxides derived from the high temperature oxidation of refractory carbides seem to suggest the role of mixed oxides to have a better protective effect on against oxidation. Testing ZrC-20TaSi2 under high temperature oxidation saw a mixture of TaZr2.75O8 oxide formed a thick dense scale including isolated porosity, which was able to halt oxidation into the underlying matrix layer [36]. Therefore, the (Zr,Ti)O2 probably functions similarly by forming a tough, sealed outer layer protecting the porous oxide layers of ZrO2 and TiO2 underneath. The mixed oxide layer lowers the diffusion rate of oxygen [37] and the oxygen partial pressure in the inner parts of the composite is depleted until it becomes too low for active oxidation of the matrix.  4. Conclusions  ZrC-20 wt% Mo and ZrC-20 wt% TiC composites synthesised by SPS, along with ZrC benchmark sample also synthesised by SPS, underwent oxidation in a static air furnace between 600 and 1200 °C in order to characterise their oxidation performance. Since ZrC oxidation is welldocumented in literature, the kinetics proceeded as would be expected with ZrO2 formation and from 600 °C and spalling of oxide flakes from 1000 to 1200 °C. Oxidation of ZrC-Mo occurs with the formation of MoO2 forming a brief interlayer between the matrix and other oxide layers, namely ZrO2. At lower temperatures, it is suspected that this MoO2 layer stabilises the oxide layer, but eventually further oxidises to MoO3. MoO3 volatises aggressively beyond 800 °C driving the weight loss of ZrC-Mo and eventual catastrophic process called pesting, where the material is reduce to fine powder. The remaining substance leaves just ZrO2 at 1200 °C on the surface of the sample. The oxidation of ZrC-TiC proceeds more gradually resisting pesting spalling even up to 1200 °C. Observation of ZrC-TiC at room  or  250  temperature showed the yellow oxide layer as a plate detaching from the matrix. It is unclear whether the loose adhesion of the oxide layer to the matrix was due to thermal stresses during cooling or coefficient of thermal expansion mismatch between oxide and matrix. The ZrC-TiC follows a complicated pathway since ZrC-TiC exists as a mixed carbide solid solution after sintering. The oxidation reaction is derived from forming mixed oxides where the species (Zr,Ti)O2 is formed, along with the standard ZrO2 and TiO2 groups. ZrC-TiC is able to resist oxidation damage better than ZrC-Mo composite or native ZrC; however, data on the (Zr,Ti)O2 species is not well-known, but the it is suspected to act as a protective barrier inhibiting further oxidation into the matrix.  Acknowledgments  This work was supported by the Estonian Research Council (PUT1063, 2015/16), The Tallinn University of Technology young researcher basic financing (B56, 2015/16), and the Institutional Research Funding (IUT 19-29, 2015/16) of the Estonian Ministry of Education and Research. The author would also like to thank Mr. Rainer Traksmaa for XRD analysis and Dr. Mart Viljus for SEM imaging.  References  [3]  [4]  [7]  [8]  [11]  [1] R. Savino, Fumo M. De Stefano, D. Paterna, M. Serpico, Aerothermodynamic study of UHTC-based thermal protection systems, Aerosp. Sci. Technol. 9 (2005) 151-160. [2] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, William Andres Publishing/Noyes, Westwood, NJ, 1996. S.E. Landwehr, G.E. 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Sciti, M. Balat-Pichelin, L. Charpentier, Zirconium carbide doped with tantalum silicide: microstructure, mechanical properties and high temperature oxidation, Mater. Chem. Phys. 143 (2013) 407-415. [31] M. Gherrab, V. Garnier, S. Gavarini, N. Millard-Pinard, S. Cardinal, Oxidation behavior of nano-scaled and micron-scaled TiC powders under air, Int. J. Refract. Met. Hard Mater. 41 (2013) 590-596. [32] K. Thamaphat, P. Limsuwan, B. Ngotawornchai, Phase characterization of TiO2 powder by XRD and TEM, J. Nat. Sci. 42 (2008) 357-361. J.M. Calatayud, P. Pardo, J. Alarcón, V-containing ZrO2 inorganic yellow nanopigments, J. R. Soc. Chem. 5 (2015) 58669-58678. [34] D. Gao, Y. Zhang, J. Fu, C. Xu, Y. Song, X. Shi, Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: formation of zircon, Corros. Sci. 52 (2010) 3297-3303. [35] D.W. Lipke, S.V. Ushakov, A. Navrotsky, W.P. Hoffman, Ultra-high temperature oxidation of a hafnium carbide-based solid solution ceramic composite, Corros. Sci. 80 (2014) 402-407. L. Charpentier, M. Balat-Pichelin, D. Sciti, L. Silvestroni, High temperature oxidation of Zrand Hf-carbides: influence of matrix and sintering additive, J. Eur. Ceram. Soc. 33 (2013) 2867-2878. F.J. Keneshea, D.L. Douglass, The diffusion of oxygen in zirconia as a functon of oxygen pressure, Oxid. Met. 3 (1971) 1-14.  [36]  [37]  Birgit Maaten did her BSc and MSc in the field of applied chemistry and biotechnology with a focus on organic chemistry and catalysis at Tallinn University of Technology. Currently doing a PhD in Thermal Engineering, with a focus on pyrolysis kinetics and thermal analysis.  Maksim Antonov graduated from Tallinn University of Technology with a PhD and has worked in the Department of Material Engineering specialising in the tribology and wear of materials. He possesses up to 80 publications to his name.  Irina Hussainova Professor at Tallinn University Technology with a broad specialisation in material engineering, with particular empathises on nanomaterials in fundamental and commercial research. She possesses over 150 publications and gives invaluable support to her many PhD students.  Der-Liang Yung is a graduate of Biochemistry BSc from the University of Simon Fraser, but later moved to Estonia to pursue a MSc Engineering in Material Science and Sustainable Energy. He is a PhD graduate in the field of Materials Engineering at Tallinn University of Technology.  251  \\x0c']"
},{
  "_id": 172,
  "PDF": "Oxidation of ultra-high temperature transition metal diboride ceramics.pdf",
  "Text": "['International Materials Reviews  ISSN: 0950-6608 (Print) 1743-2804 (Online) Journal homepage: https://www.tandfonline.com/loi/yimr20  Oxidation of ultra-high temperature transition metal diboride ceramics  W G Fahrenholtz & G E Hilmas  To cite this article: W G Fahrenholtz & G E Hilmas (2012) Oxidation of ultra-high temperature transition metal diboride ceramics, International Materials Reviews, 57:1, 61-72, DOI: 10.1179/1743280411Y.0000000012  To link to this article:  https://doi.org/10.1179/1743280411Y.0000000012  Published online: 12 Nov 2013.  Submit your article to this journal   Article views: 885  View related articles   Citing articles: 109 View citing articles   Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=yimr20  \\x0c', 'Oxidation of ultra-high temperature transition  metal diboride ceramics  W. G. Fahrenholtz* and G. E. Hilmas  The oxidation behaviour of  transition metal diboride ceramics is reviewed with emphasis on the  performance  of  zirconium diboride  and hafnium diboride.  First,  the  oxidation behaviour  of  nominally pure diborides is discussed,  focusing on the transition to linear mass gain kinetics at  temperatures above y1100uC. Next, the use of SiC and other additives that produce silica based  scales when oxidised is  reviewed. These additives  improve oxidation protection due to the  formation/stability of  the outer  layer of borosilicate glass that acts as a barrier  to diffusion of  oxygen to the substrate. However, elevated temperatures (.1650uC) and/or  the combination of  aerodynamic flow, high heat  flux and reactive atmosphere associated with hypersonic flight  destabilises the outer oxide and decreases oxidation protection. Other additives that affect  the  composition and structure of the crystalline oxide scale without forming an outer glassy layer are a  promising approach to improving oxidation behaviour of diborides. These additives  require  further  research to understand the mechanisms of  improved protection and further optimise the  protective behaviour. While the oxidation of ultra-high temperature diborides has been studied for  many years, several possible areas for  future research are identified.  Keywords: Zirconium diboride, Hafnium diboride, Oxidation, Review  Introduction  Ultra-high temperature ceramics (UHTCs) are a family  of materials that include a number of borides, carbides and nitrides of early transition metals.1 One criterion that  has been used to define this family of materials is a melting temperature of 3000uC or higher, although other criteria such as the ability for continuous use at temperatures above 1600uC can also be used. From this broader family of materials, the refractory diborides, and  more  specifically ZrB2  and HfB2,  exhibit  an unusual  combination of ceramic-like strength (500 MPa or higher room temperature)2,3 and elastic modulus (y500 at room temperature)4 with metal-like elec(y107 S m21) (.60 W m21 K21) trical and thermal conductivities.5 Some of the relevant physical properties of ZrB2 and HfB2 are summarised in Table 1.6,7 Based on their combination of properties, ZrB2 and HfB2 are  at  GPa  candidates for applications ranging from high tempera ture  electrodes,  cutting  tools  and molten metal conlayers.8-13 While  tainment  to microelectronic  buffer  transition metal diboride compounds have been studied for over 100 years,14 it was not until  the 1940s and 50s  that the first significant processing-structure-property studies were reported.15,16 Through the 1960s and early  1970s, ZrB2 and HfB2 emerged as candidates for a variety  of aerospace applications, which resulted in significant  research efforts in both the US and the former Soviet Union.17-19 During the past 15 years,  these compounds  have been the subject of renewed attention for aerospace  applications related to hypersonic flight, rocket propulre-entry.20-23  sion  and  atmospheric  In  particular,  the  refractory diborides possess a combination of strength at  elevated  temperatures  and  thermal  conductivity  that  provides improved thermal shock performance in high heat flux conditions compared to other ceramics.24-26 As  a result of their desirable properties, ZrB2 and HfB2 are  candidates for components on future hypersonic vehicles  that undergo intense aerodynamic heating such as leading inlets.27-30  and trailing edges or engine cowl  The US  Space  Shuttle Orbiter  has  been  the  only  reusable  atmospheric  re-entry  vehicle  for  the  past  30 years. The Shuttle  employs a blunt  edge design to  mitigate the intense heat associated with atmospheric re-entry.31 As shown schematically in Fig. 1a, the blunt  surfaces produce a shock wave ahead of the vehicle that  deflects  some of  the heat  away  from the  surface by  transferring much of the kinetic energy to the air behind  the vehicle rather than the leading edges or other vehicle surfaces.32 The combination of  re-entry trajectory and  vehicle design limits the maximum surface temperatures to y1650uC, but  on the nose  cap and leading edges  reduces  the maneuverability  and  cross-range  of  the  Shuttle.  In contrast, hypersonic vehicles which employ  sharp leading edge designs increase maneuverability and cross-range.33  Increased maneuverability requires  lami nar flow across  the  control  surfaces, which,  in turn,  necessitates the use of sharp leading edges. Although the  Department of Materials Science and Engineering, Missouri University of  Science and Technology, Rolla, MO 65409, USA  *Corresponding author, email billf@mst.edu  ß 2 0 1 2 I n s t i t u t e o f M a t e r i a l s , M i n e r a l s a n d M i n i n g a n d ASM I n t e r n a t i o n a l P u b l i s h e d b y M a n e y f o r t h e I n s t i t u t e a n d A SM I n t e r n a t i o n a l  DO I 10 . 11 79 /1 74 32 8 04 11 Y . 00 00 00 00 12  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  6 1  \\x0c', 'maximum surface temperatures depend on factors including speed and radius of curvature,34 temperatures in excess of 2000uC are predicted for sharp leading edges of most  hypersonic vehicles due to the high heat flux impinging directly on the sharp tip.35 As a result of the heat flux at  the sharp tip, heat must be conducted away from the tip  and through the  leading edge  so that  it  can either be  dissipated by reradiation from cooler surfaces away from  the leading edges (Fig. 1b) or transferred internally to a system.36  more  complex,  active  cooling  The  surface  temperatures encountered by sharp leading edges or as  other potential applications mean that oxidation resistance is a critical property of the refractory diborides.37  Historical studies concluded that the relative oxidation  resistances of ZrB2 and HfB2 were superior to those of other transition metal diborides.38 Hence, studies focused  on aerospace applications have  concentrated on these  materials as will this review. The purpose of this paper is  to critically review historical and recent research related  to the oxidation behaviour of ultra-high temperature  diboride ceramics with emphasis on ZrB2 and HfB2.  Oxidation of transition metal diborides  Zirconium and  hafnium diborides  undergo  stoichioand (2).39  metric oxidation according  to reactions (1)  The expressions for the change in standard state Gibbs’ free energy with reaction (nGurxn) were calculated for the temperature range from room temperature (y25uC to y2000uC (2275 K) using data from the tables40 standard reference (reaction (1)) dynamic software (reaction (2)).41  or 300 K)  and thermo ZrB2(cr)z  5  2  O2(g)?  ZrO2(cr)zB2O3 (1) DGo rxn  ~{1999500z374:4T (J)  (1)  Hf B2(cr)z  5  2  O2(g)?  Hf O2(cr)zB2O3 (1) DGo rxn  ~{2000460z375:7T (J) (2)  The reaction products, B2O3 and either ZrO2 or HfO2,  show limited/no mutual solubility, so the scales contain  two distinct phases. As with metals and other non-oxide  ceramics,  oxidation  of  these materials  is  favourable  across a wide range of temperatures that includes from room temperature to .2000uC. The implication is that differences in oxidation behaviour depend on kinetic  factors since a strong driving force exists across a wide  temperature range. As such, the oxidation behaviour of  the phase-pure diborides  is divided into two regimes  based on whether the oxide that  forms is protective (at  ‘low’  temperatures) or not (at ‘high’  temperatures).  Both  ZrB2  and HfB2  exhibit mass  gain  kinetics  consistent with diffusion limited processes  in the  low  temperature  regime. The  upper  limit  of  this  regime  depends on factors  such as  external pressure, oxygen  partial  pressure  and  gas  flowrate, but is generally and 1200uC in static air. Below the transition temperature, a protective oxide  considered to be between 1100  scale forms on the surface of ZrB2 and HfB2 and both  ceramics  show parabolic  trends  for mass  gain,  scale  thickness and oxygen consumption time.39,42 Historical43 and current  as  a  function  of  (Fig. 2a) analyses of  1  Notional  representations of  leading edges of hypersonic  aerospace  vehicles with a blunt  and b sharp leading edges  showing the effects that  lead to heating by convective flow to the surface qconv,  radiation through the boundary layer  qrad,  chemical  reactions  in  the  boundary  layer qchem and surface  catalysis qcatal with  heat dissipation  by  conduction  away from the surface qcond and reradiation qrerad  Table 1  Overview of physical properties of ZrB2 and HfB2  ZrB2  HfB2  Reference  Melting temperature/uC Crystal structure  3245  3380  6  Hexagonal AlB2  Hexagonal AlB2  7  Space group Theoretical density/g cm23  C6/mmm 6.1  C6/mmm 11.2  8  4  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  6 2  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  \\x0c', 'cross-sections of oxidised specimens  reveal a two layer  oxide scale that consists of an outer layer of glassy B2O3  and  an  inner  layer  that  contains  porous ZrO2 with  the pores filled by glassy B2O3. More recently, Parthasarathy et al.44-46 developed an oxidation model  for TiB2, ZrB2 and HfB2 showing that the oxidation rate  is limited by the diffusion of oxygen through B2O3 (i.e.  transport of oxygen through ZrO2 is negligible). Based  on the combination of historical and recent experimental  and modelling results, diborides exhibit passive oxida tion behaviour with the formation of a protective oxide  scale in the low temperature regime. In the high temperature regime (i.e. above y1200uC), the oxidation behaviour of the diborides changes.42,47-49 (previous43,50  Microstructural  analysis  and  Fig. 2b)  reveals  that  loss  of  protection  for  ZrB2  coincides  with  evaporation  of  B2O3  from  the  oxide  scale.  Thermodynamic models diagrams37,51 or kinetic models, such as the one proposed by Parthasarathy et al.,44 support the evapora that  employ  either  volatility  tion  of  B2O3  as  the  cause  of  the  transition.  Thermodynamic models  predict  vapour  pressures  of  the various gaseous boron oxides that form as a function  of  external  conditions  such  as  temperature,  oxygen  partial pressure, etc. As shown in Fig. 3, B2O3 (g) is the  predominant vapour species formed by evaporation of B2O3 in air at 1500uC. Although B2O3 volatilises over a  wide range of conditions, changes in the partial pressure  of  oxygen  in  the  external  atmosphere  (shown  on  diagram)  or  temperature  at which  oxidation  occurs  (shown  in  volatility  diagrams  in Ref. 51)  affect  the  predominant species in the vapour phase. HfB2 exhibits  the best oxidation resistance of  the diborides over  this  temperature range because the oxide layer that remains  after B2O3  evaporation  has  a more  equiaxed micro structure, which gives it greater transport.39 Although ZrB2  resistance  to oxygen  is  inferior  to HfB2, both  have significantly better oxidation protection than other diborides such as TiB2, TaB2 and NbB2.1  Direct comparisons of historical and recent oxidation  results  for nominally pure diborides  is difficult due to  lack  of  convention  in  reporting  results. Whereas  historical studies have used a combination of mass gain,  scale  thickness,  parabolic  rate  constant  and  oxygen  uptake  as  a  function of  temperature  and time, more  recent  studies  have  focused  on  thermal  gravimetric  3  a vapour pressure of various B-O species as function of oxygen partial pressure at 1500uC and b ZrB2 volatility dia gram based on calculations described in Ref. 51  2  Cross-section images of oxide scale on nominally pure  ZrB2  oxidised  in  air  at  a  900uC  for  8 h  (oxide  layer  thickness, y10 mm)  and  b  1500uC for  2 h  (oxide  layer  thickness, y400 mm)  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  6 3  \\x0c', 'analysis  (TGA)  to measure mass gain as a function of  temperature  and/or  time.  In  addition,  differences  in  oxidation temperature and time also complicate direct  comparison, but Table 2 provides an overview of some reports.52 Where direct  historical  and recent  compar isons are possible  for nominally phase-pure diborides,  the trends are consistent, but the quantitative values do not agree. For example, Tripp and Graham49 reported a 3?3 mg cm22 mass gain of for pure ZrB2 heated to 1300uC for 2 h, which is in the linear kinetic region. For comparison, Opeka al.53 reported a mass gain of 9?8 mg cm22  et  for  the  same  oxidation  conditions. No  obvious differences  in density,  composition or micro structure can be identified as the cause of the difference  in mass gain. Based on the model of Parthasarathy et al.,44 small changes in the fraction of porosity in the  ZrO2 have  a  significant  effect on the oxidation rate.  Trace impurities in the diborides are one potential cause  of the differences in oxidation behaviour, an idea that is  discussed in more detail  in the sections that  follow.  As mentioned above,  research related to aerospace  applications has  focused on ZrB2 and HfB2. However,  because of  the  technological  importance of TiB2  in a  number  of  applications,  several studies, Parthasarathy’s modelling reports,44 have examined the oxidation resistance of TiB2.50,54 Only the historical  including  studies performed by Manlabs examined the oxidation  resistance of other nominally phase-pure diborides detail,55  in  although  one recent report discussed oxidation resistance of TaB2.56 By TGA, nominally phase-pure TaB2 gained y20 mg cm22 when heated to to y6?5 mg cm22 1500uC compared for nominally heated under the same  the  phase-pure ZrB2  conditions.  This  study  did  not  provide  new  insight  into  the  behaviour, but served to reinforce previous conclusions  about  the  increased relative oxidation rates of other  diborides.  Furnace oxidation studies,  such as  those described  above, do not reproduce conditions that are representa tive of  extreme  environments  such as  the aerothermal  heating encountered during hypersonic flight. While the  conclusions of  these studies provide insight  into oxida tion mechanisms  and may  be  useful  for  screening  candidate materials, hypersonic flight produces higher  heat fluxes, dissociated gaseous species and gas flowrates  that cannot laboratory furnaces.1 Specialised facilities such as arc heaters,57 plasma wind tunnels58 and inductively coupled plasma facilities59 have been developed to test materials in more  be  duplicated  in  typical  realistic  environments.  Only  a  limited  number  of  published  papers  have  reported  the  behaviour  of  nominally phase-pure diborides  in environments  rele vant to hypersonic flight. For surface temperatures reaching up to y2500uC (no heat flux reported), oxide reached y250 mm in 1800 s scale thicknesses on ZrB2 (30 min).60 For comparison, the same material had an oxide scale y2?5 mm (2500 mm) thick after furnace oxidation at y2000uC for 1800 s (30 min). The harsher conditions of the arc heater may have resulted in a  thinner scale due to loss of material by evaporation and/  or flowoff of  the surface during testing. However,  the  limited number of experimental studies do not allow for  strong conclusions to be drawn about  the behaviour of  the material. Further,  it  could be  concluded that  the  more  severe  conditions  encountered during arc heater  testing may promote  the  formation of an oxide  scale  (ZrO2 or HfO2) having a higher density that  is more  protective  than the  scale  formed in static  laboratory  furnace tests.  Oxidation of diboride ceramics containing SiC  The  relatively poor oxidation resistance of nominally temperatures above y1200uC to investigate a number of  phase-pure diborides at  motivated  researchers  approaches to improving oxidation resistance including  solid solution additions,  synthesising ternary diboride phases.61 The most  compositions  and adding  second  promising approach was found to be the addition of SiC  as a second phase, which reduced the thickness of  the  oxide  scale  across  a  wide  temperature  range  when  compared to either (Fig. 4).61,62 The  a  pure  diboride  or  pure  SiC  improved  oxidation  resistance was  attributed to the formation of a stable borosilicate glass layer on the surface of the oxidised ceramics.43 Since this  pioneering work,  research on oxidation resistant ZrB2  and HfB2 based compositions with silica scale forming  additions has  resulted in a significantly larger body of  work on SiC containing materials  than other  systems.  The rest of  this section discusses recent progress on the  most  common approach for  improving  the oxidation  behaviour of diboride ceramics, which is the addition of  SiC.  Several  groups  have  studied  the  effects  of  SiC  additions on the behaviour of diborides when exposed  to air at elevated temperatures. The reported behaviour  is  summarised in Table 3 and discussed in more detail  Table 2  Summary of historical and recent oxidation results for nominally pure ZrB2 and HfB2 ceramics  Material  Oxidation conditions  Weight gain/mg cm22  Parabolic rate constant KP/kg2 m24 s21  Reference  ZrB2  1000uC, 4 h 1500uC, 4 h 1200uC, 2 h 1300uC, 2 h 1400uC, 2 h 1522uC, 60 min 1522uC, 58 min 1166uC, 1 h 1256uC, 1 h 1025uC 1400uC, 2 h 1500uC, 2 h 1600uC, 2 h  1.1 6.6 2.2 2.8 5.6  …  43  ZrB2  …  49  HfB2  …  561026 161024 761027 961027 961024  39  ZrB2  ZrB2  4.8 5.7  47  ZrB2  … 18.7  50  ZrB2  …  53  ZrB2  11  …  52  14  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  6 4  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  \\x0c', 'below.  Below y600uC, ceramics to air or oxygen results in no significant mass gain,63 which indicates  exposure  of  diboride  based  that oxygen transport  through  the native oxide layer  is negligible in this temperature to y1100uC, 1-2 mg cm22  regime. As  the  temperature  increases  cumulative mass  gains on the order of by TGA,64,65  have  been  observed  indicating  oxygen  transport  through the oxide layer becomes more rapid. than y1100uC, showed that the oxidation behaviour of  For  temperatures  less  mass  gain  measurements  SiC containing diborides was not affected by the SiC  additions. Examination of  the oxide scale revealed that  ZrB2 or HfB2 oxidised preferentially at  these tempera tures, which produced an oxide scale containing B2O3, ZrO2 or HfO2, and leaving the SiC mainly unoxidised.66 to 1100uC and above, B2O3 evaporation becomes significant and SiC begins to  As  the  temperature  increases  oxidise. The  formation  of  borosilicate  glass  in  this  temperature  regime  results  in  parabolic mass  gain  kinetics with parabolic 1028-1026 kg2 m24 s21, as reported authors.66-68 Hence, oxidation at  rate  constants  in the  range of  by  several  temperatures accessi ble in typical  laboratory furnaces with air atmospheres (i.e. up to y1600uC) with the formation of a protective oxide layer composed  or TGA equipment  is consistent  of borosilicate  glass, a conclusion that has also been computational modeling.69,70 After  supported by  fur nace  oxidation  at  temperatures  between  1100  and  1600uC, the formation of a SiC depleted region has been observed by some researchers (Fig. 5a),38,43,64,66,67 (Fig. 5b).63,71,72  but not by others  In this  temperature  regime,  the oxide  scale  consists of  an outer  layer of  borosilicate glass, a layer of mixed ZrO2 and SiO2, a  partially oxidised layer, and the underlying unoxidised  ceramic. This  structure was  shown most distinctly in 1627uC  ZrB2-SiC that underwent (Fig. 6).66  cyclic oxidation at  The  depletion  of  SiC  from the  partially  oxidised layer has been attributed to active oxidation  of SiC due to the oxygen activity gradient through the oxide.73 Without  outer  layer  of  dense  glassy  SiC  depletion,  the  partially  oxidised  layer  can  consist  of  either ZrO2 or a graded structure  consisting of ZrO2  closer  to  the  outer  surface  and ZrB2  closer  to  the  unoxidised material. For SiC depletion,  either  type of  partially oxidised layer would then also contain voids  corresponding to the places where SiC particles were  present  in  the  original material. No  study  has  yet  addressed  the  reasons  that  SiC depletion  has  been  observed in some materials, but not others; however, a  combination  of  compositional  and  microstructural  factors  likely control  the microstructure  that develops  in  the  oxide  scale  (i.e.  SiC depletion  or  not). For  example, assuming that SiC depletion relies on active  oxidation of SiC, then depletion could only occur when  the  SiC  particles  formed  an  interconnected  three dimensional network.  Far  fewer  studies  have  examined  the  oxidation  behaviour of diboride materials at temperatures beyond y1600uC. Research from the 1960s and 70s evaluated oxidation behaviour up to y2200uC,61 but more recent to y1600uC. studies have typically been limited  Thermodynamic  analysis predicted loss of protection temperatures above 1600uC due to  of the oxide scale at  gas pressure beneath the scale causing bubbles/voids in the scale,37 which was later confirmed by in situ observations of oxidising ZrB2-SiC.74 Although bubble  formation was predicted to occur at temperatures as low as 1450uC for oxidation times of .6 h,  the incubation  period required before bubble nucleation was much lower at 1600uC and above, decreasing from y1 h at 1600uC to nearly 1650uC.74 Pushing  instantaneous 1800uC  at  temperatures  to  and  above  causes  bubble  formation and eventual loss of the glassy oxide along the crystalline oxide and voids.75-77  with coarsening of  In  this  temperature  regime,  oxidation  behaviour  is  Table 3  Summary of oxidation behaviour of SiC containing ZrB2 and HfB2 ceramics  Temperature range  Scale  Oxidation behaviour  ,600uC 600-1100uC  Native oxide ZrO2 or HfO2zB2O3 SiC not oxidised  Little,  if any mass gain/scale growth  Parabolic kinetics due to presence  of a dense glassy oxide layer (B2O3)  1100-1300uC  Borosilicate glass outer layer ZrO2 or HfO2zSiO2 layer  B2O3 evaporation and SiC oxidation  Paralinear kinetics: evaporation of B2O3  plus protection from borosilicate glass  1300-1600uC  Borosilicate glass outer layer  B2O3 evaporation, ZrO2 transport,  and SiC oxidation  ZrO2 or HfO2zSiO2 layer  Parabolic kinetics due to dense  borosilicate glass layer  Partially oxidised layer  SiC depletion observed in some cases  .1600uC  Gas pressure at scale/boride  interface and scale evaporation  Bubbles in scale, presumably CO  Evaporation of silica from scale  The ZrO2 scale can be protective or not  4  Comparison  of mass  gain  as  function  of  temperature  for 1 h oxidation in air  for HfB2, SiC and HfB2-SiC (re printed with  permission  from Ref. 61,  copyright  1968,  The Metallurgical mil<25 mm  Society  of  AIME):  note  that  1  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  6 5  \\x0c', 'Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  5  Scanning electron micrographs of oxidised diboride  ceramics  containing SiC,  a ZrB2-SiC oxidised at  1500uC showing  partial SiC depletion (reprinted with permission from Ref. 64,  copyright  2007, Elsevier)  and b the outer  and intermedi ate layers on HfB2-SiC oxidised at 1350uC showing no SiC depletion (reprinted with permission from Ref. 63, copyright  2003, The Electrochemical Society)  determined by the microstructure of the crystalline oxide (ZrO2 or HfO2) present on the surface.75  to discoveries related to liquid convection and oxide transport during oxidation.79,80 Notably,  particle  this  In addition to furnace oxidation and TGA studies  technique has  led to the identification of zirconia rich  that have been used to determine oxide layer thicknesses  nodules  that  grow up  and  through  the  liquid  oxide  and mass gains as functions of time and/or temperature,  (Fig. 7). The presence of  these nodules has provided  recent progress has been made in the use of  real  time  insight  into  potential  causes  for  the  variability  in  observations of growing oxide layers. Because diborides  are electrically conductive, direct heating of specimens is  possible using conventional furnace power supply and systems.74,78 Direct heating allows observation  control  thicknesses of the outer glassy oxide and underlying partially oxidised regions noted by some authors.68,81,82  The  external atmosphere has a strong effect on the  oxidation  behaviour  of  diboride  based  ceramics.  of the specimen surface during oxidation, which has led  Because of proposed applications in propulsion systems,  6  The  layered structure  that developed on ZrB2-SiC oxidised by heating to 1627uC for  10 cycles. The oxide  scale  con sists of an outer  layer of SiO2, a layer of ZrO2 and SiO2, and a layer of SiC depleted ZrB2 on top of unreacted ZrB2-  SiC (reprinted from Ref. 66 with permission, copyright 2001, The American Ceramic Society; all  rights reserved)  6 6  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  \\x0c', 'the effect of water vapour was examined on the oxidation behaviour of HfB2 and ZrB2 ceramics.83 The  oxidation  rates  of  the UHTCs were  comparable  in  stagnant air and an atmosphere of 90% water vapour  and  10%  oxygen.  However,  the  UHTCs  showed  significant degradation in flowing water vapour due to silica volatilisation.83  In addition,  the effect of oxygen  partial pressure has been evaluated in historical and recent studies.61,75,84 At oxygen partial pressures ranging from pure oxygen (1 atm or 101 325 Pa) to y20 Pa, the scale thickness and mass gains appear comparable.68,75  However, oxidation at an oxygen partial pressure of y10210 Pa at 1500uC resulted in active oxidation of SiC  along with conversion of the ZrB2 to ZrO2 with the loss of B2O3 by evaporation.84 The latter experiment verified thermodynamic predictions regarding the mechanism of  SiC depletion by  active oxidation that was proposed  based on experimental observations and thermodynamic calculations.66,73  Other additives have been used in combination with  SiC to  further  improve  the  oxidation  resistance  of  refractory diborides. Opila investigated the addition of  TaSi2 to ZrB2-SiC and HfB2-SiC ceramics and found improved oxidation resistance in ZrB2-SiC at 1627uC, but degraded performance at higher temperatures.85 The  addition of TaSi2 did not affect the oxidation behaviour of HfB2-SiC.85 Talmy et al.37  studied the addition of  10 mol.-% CrB2, TiB2, NbB2, VB2 or TaB2 to ZrB2-SiC  and showed that  the additives  increased resistance  to  oxidation (Fig. 8). The authors reported that  the outer  glassy layer underwent phase separation due to the high  cation field strength of  the  transition metal additives,  which improved the oxidation resistance of the resulting  ceramics by increasing the viscosity and decreasing the  oxygen diffusivity through the oxide scale that formed ceramic.37 More  on the  surface of  the diboride  recent  studies have revealed that addition of TaB2 and/or TaSi2  to ZrB2-SiC modified the morphology of ZrO2 in the oxide scale produced at temperatures up to 1500uC.86,87 (e.g. 3?3 mol.-%), For additions of a few mole per cent  the oxide scale had grains with a more equiaxed shape  which increased retention of the borosilicate glass and improved oxidation resistance.86 For higher contents of  Ta  based  compounds,  the  additions  appeared  to  promote  coarsening  of  the  crystalline  phases  in  the  oxide  scale and the  formation of ZrO2 dendrites  that  penetrated  the  outer glassy layer, both of which protection.87 Complementing  degraded  oxidation  the  use of Ta based additives,  the oxidation behaviour of  TaB2-SiC has also been examined, with TGA studies  revealing that TaB2 containing 20 vol.-%SiC gained y3?0 mg cm22 when heated to 1450uC,88 compared to y2?5 mg cm22  for  ZrB2  containing 20 vol.-%SiC conditions.65 LaB6  heated under nominally  the  same  was also identified as a beneficial additive to ZrB2-SiC  based on its ability to act as a modifier for the ZrO2, which stabilised the formation of the tetragonal phase.89  Finally, the use of nitride additives as sintering aids has  been shown to lead to rupture of the oxide scale and loss  of protective behaviour 1400uC.72 These  at  temperatures  as  low  as  studies have demonstrated that addi tives  affect  the  oxidation  behaviour  of  diboride-SiC  ceramics by altering the structure of the crystalline oxide  and/or  glassy  layers  in  the  oxide  scale. However,  additional research is needed to determine the mechan isms by which additives affect the structure of the oxide  7  Images of ZrB2-SiC after oxidation at 1550uC showing a lighter coloured ZrO2  rich nodules protruding through the dar ker glassy  surface oxide  layer  (reprinted from Ref. 74 with permission,  copyright  2010, The  Journal of  the European  Ceramic Society)  and b the  cross-section of  a  lighter  coloured ZrO2  rich nodule  extending through the outer glassy  oxide layer  (reprinted from Ref. 79 with permission, copyright 2007, The Journal of  the American Ceramic Society)  8  Comparison  of mass  gains  for  nominally  pure  ZrB2-  SiC with  ZrB2-SiC ceramics  containing  other  transition  metal  diboride  additives  (reprinted  from Ref. 37  with  permission,  copyright  2004,  The  Journal  of  the  Materials Science)  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  6 7  \\x0c', 'scale along with the impact on transport of oxygen and/  or gaseous oxidation products through the scale.  Arc  heater  or  similar  testing  of  SiC  containing  diborides has been much more  extensive  than for  the  nominally phase-pure diborides. Testing has  revealed  two regimes of behaviour:  (i)  a low temperature regime in which a glassy oxide  is maintained on the specimen surface  (ii)  a high temperature regime in which the specimen  surface is primarily crystalline oxide.  These regimes are discussed in more detail below.  The low temperature regime is defined by surface temperatures less than y1600uC. Generally, this condition results for cold wall heat fluxes that are less than y300 W cm22. However, two regimes depends on  the  transition  between  the  test  conditions  such  as  stagnation pressure,  gas flowrate,  composition of  the  atmosphere, surface catalytic efficiency of the specimen,  sample  geometry,  etc. Therefore,  the  real  heat  flux  experienced  by  the  specimen  during  testing  is  not  known,  leaving  the  surface  temperature  as  the  only  measureable parameter to denote the transition. When a  glassy surface layer  is present, diboride based UHTCs  exhibit behaviour consistent with non-catalytic surfaces,  which reduces surface heating due to recombination of species.90 Analysis of  dissociated gas  time dependent  evolution  of  gaseous  species  from ZrB2-SiC  in  a  plasmatron revealed that boron based species volatilised  during  the initial stages present after y150 s y120 W cm22.91 The ceramics showing similar evolution of  of  exposure,  but were  not  of  exposure  at  a  heat  flux  of  results point  to diboride based  structure during  arc jet exposure as in furnace oxidation studies in which  ZrB2 oxidises at  lower  temperatures producing a B2O3  rich  scale,  which  volatilises  at  higher temperatures SiO2.64 thermodynamic calculations, the presence of monatomic  leaving  a  scale  that  is  primarily  Based  on  oxygen  appears  to  increase  the  driving  force  for  oxidation of SiC, but does not have a significant effect rich surface oxide.92  on volatility of  the resulting SiO2  The presence of small volume fractions of additions such  as AlN does not appear to have an adverse effect on the  arc jet performance of diboride ceramics when the glassy present.93 However,  surface  layer  is  comparison  of  HfB2-SiC compositions produced by conventional hot  pressing and spark plasma sintering revealed that  the  size and distribution of the SiC particle affected arc heater performance.94 For  inclusions  ceramics with  equal volume fractions, SiC particles that were smaller  and more uniformly distributed led to the formation of  thinner glassy surface  layers  for  similar  testing condi tions. The conclusion was  that  the distribution of SiC  particles  affected  the  rate  at which  the  glassy  layer  formed, with finer SiC particles distributed uniformly  producing a thin, but protective layer more quickly and,  therefore, improving the overall oxidation performance the ceramic.94 Based on the studies discussed above,  of  behaviour of diboride-SiC ceramics in the low tempera ture  regime  is  controlled by the presence of a  glassy  surface layer, which is similar to the behaviour reported  for furnace oxidation y1600uC. Hence, oxidation may be applicable to the arc jet behaviour of diboride based UHTCs at heat fluxes of y300 W cm22 or lower.  studies  at  temperatures  of  conclusions  drawn  from furnace  At higher heat fluxes, those temperatures above 1600uC,  that  produce  surface  the behaviour of diboride  based ceramics changes (Fig. 9). At heat fluxes of more than y300 W cm22,  the  glassy  surface  layers  are  removed, which leaves a surface covered with crystalline  oxide (i.e. ZrO2  for ZrB2 based ceramics and HfO2  for  HfB2 based ceramics). In addition, the formation of a SiC  depleted region is  reported after  testing at  these condi tions, but  test articles appear  to maintain their dimen sional stability (i.e. no significant recession or deformation).95 Testing at a heat flux of y350 W cm22  in the same facility produced surface temperatures of y1800uC after 600 s on ZrB2-SiC,65 but y2360uC after 600 s on HfB2-SiC.57 The higher surface temperature for  HfB2-SiC was attributed to a higher catalycity of HfO2  and the lower thermal conductivity of the HfO2 surface oxide compared to HfB2,57 although the authors pointed  to the need for more extensive and comprehensive testing  as their conclusions were based on a limited number of  specimens. Unlike  testing  at  lower heat fluxes, which  showed little dependence on the presence of additives, testing of ZrB2-SiC containing either TaSi2 85 or AlN,93  showed that significant mass loss and recession occurred  at conditions  that did not  result  in either mass  loss or  recession for ZrB2-SiC without other additives. Finally, a  coordinated effort was used to select promising candi dates using screening studies and then evaluate the effect  of fabrication procedures, performance.58,90,96-99  shape and other  factors on  The  analysis  showed  that  at  moderate heat fluxes that produced surface temperatures  9  Composite  image of HfB2-SiC ceramic of y350 W cm22  after  arc heater  testing  at  heat  flux  showing  crystal line ZrO2 surface layer, SiC depleted layer and underly ing  ceramic  (reprinted  from Ref. 57  with  permission,  copyright 2004, The Journal of  the Materials Science)  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  6 8  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  \\x0c', 'of y1800uC, both sharp and blunt tained their shape during exposure and that the degree of SiC depletion depended on the local specimen geometry.58  test models main Going beyond .500 W cm22 led to significant mass losses and surface recession.93 Although enough testing has been reported in  the  ‘moderate’  regime  to  heat  fluxes  the open technical  literature to identify trends  in surface  temperatures, recession and shape stability, a focused and  coordinated research effort  is needed to identify funda mental aspects of composition, microstructure and proper ties that control behaviour of diboride based UHTCs in arc  heater testing.  Other approaches to improving oxidation resistance  Diboride based ceramics with SiC additions have been  researched extensively, but silicides and other Si contain ing additives also produce a glassy silica rich surface  layer during oxidation. Hence, other additives provide  improvements in oxidation protection similar to SiC. Of  particular note  is  an extensive body of work on the  processing, microstructure  and  properties  of MoSi2  containing UHTCs has been generated by Monteverde  and  co-workers  at  the  Institute  for  Science  and  Technology of Ceramics (an Italian government labora tory in Faenza,  Italy);  researchers  there add MoSi2  to  promote the pressureless sintering of ZrB2, HfB2 and carbide UHTCs.100-102 Like SiC additions, MoSi2 is  present as a second phase in the final microstructure and  produces a SiO2 rich outer oxide layer for oxidation at temperature above 1200uC.103 The glassy surface oxide protects the underlying ceramic from oxidation and kinetics.103  results  in  parabolic mass  gain  Table 4  compares mass gain and parabolic  rate  constants  for  ZrB2  ceramics  containing  either  20 vol.-%SiC or  20  vol.-%MoSi2, which shows that the rate constants were the mass gain was y6?5 mg cm22 for the containing ceramic compared to y3?0 mg cm2 for ZrB2-SiC. The higher mass gain for ZrB2-MoSi2  similar, but  MoSi2  was attributed to retention of Mo in the oxide scale as  compared to SiC, which releases carbon as CO during oxidation.103 Therefore, added mass gain may not be an  indication  of more  severe  oxidation. Based  on  these  results, silicide additions, without SiC, are a promising  alternative  that may  alleviate  issues  associated with  formation of a porous SiC depleted region as has been  observed for SiC containing diborides.  In addition to MoSi2, other examined as additives to improve oxidation resistance of  silicides have also been  diboride ceramics. As with SiC and MoSi2, additions of other silicides also produce a silica rich surface scale  when oxidised. The oxidation behaviour of HfB2-TaSi2  was found to be similar to that of diborides with MoSi2 to y1900uC.104 not alter the oxidation  additions  for  temperatures  up  The  addition  of  1 wt-%ZrSi2  did  behaviour  of  ZrB2,  but  larger  additions  results  in  parabolic mass  gain kinetics at 1200uC.105 For ZrB2 (y18 vol.-%) the mass gain 2 h at 1200uC. Talmy investigated the oxidation behaviour of ZrB2 containing  containing 15 wt-%ZrSi2 2?2 mg cm22  was  after  et  al.  several different silica scale formers including Ta5Si3, Si3N4 and TaSi2.106,107 For Ta5Si3 additions, oxidation resistance was produced for additions y4 mol.-% (15 vol.-%), which resulted in mass gains of y5 mg cm22 after 2 h at 1200uC.106 For additions in this range, the Ta5Si3 dissolved into the ZrB2 matrix to  the best  of  form  a  solid  solution. More  important  than  the  magnitude  of  the  mass  gains,  these  compositions  exhibited parabolic mass gain kinetics with the formation of a silica rich outer scale.106 On their own, Si3N4 additions of y35 vol.-% were needed to provide to y25 vol.-%SiC at oxidation protection equivalent temperatures up to 1400uC.107 However, the addition of 20 vol.-%Si3N4 with 10 mol.-% of either CrB2 or TaB2  to ZrB2 resulted in mass gains that were comparable to ZrB2-SiC at temperatures up to 1400uC.107 The addition of the transition metal borides to ZrB2-Si3N4 modified  the  composition  of  the  crystalline  oxide  scale  and  resulted in phase  separation of  the glassy outer  layer  as had been observed for the addition of transition metal borides to ZrB2-SiC by the same research group.37 The results of these studies reinforce the conclusions drawn  earlier that further research is needed to understand the  function of additives on the development of oxide scales  as well as the relative importance of  the glassy surface  scale and the underlying crystalline oxide in the overall  oxidation protection of diboride based UHTCs.  Building on the previous point,  the oxidation resis tance  of  refractory  diborides  can  be  improved with  additives  that modify  the  structure of  the  crystalline  oxide scale without forming glassy oxides. The oxidation  resistance of ZrB2 ceramics containing up to 8 mol.-%W has been investigated.52,108 The presence of W dissolved  into the ZrB2 matrix was found to reduce both mass gain  and oxide scale thickness compared to nominally pure ZrB2 at 1500 and 1600uC. For example, the mass gain after exposure to 1600uC for 2 h was y15 mg cm22 for nominally pure ZrB2, but decreased to y8 mg cm22 for ZrB2 containing 4 mol.-%WC.52 The improved oxida tion resistance was attributed to the formation of a ZrO2  scale with  a  crystalline,  equiaxed microstructure  and  retention of a combination of WO3 and B2O3 scale.108 Further, mass gain kinetics appeared to be and 1600uC.52,108 Hence, oxidation behaviour can be improved solely by manipulating the  in the  parabolic  at  1500  morphology of  the crystalline phase in the oxide scale  whereas most studies have focused on the effect of  the  outer glassy layer.  Summary and outlook  This paper has  reviewed historic and recent  reports of  the  oxidation  behaviour  of  ultra-high  temperature  Table 4  Comparison of oxidation behaviour of ZrB2 containing 20 vol.-%SiC or 20 vol.-%MoSi2  Material  Mass gain*/ mg cm22  Parabolic rate constant KP/kg2 m24 s21  Reference  ZrB2z20 vol.-%SiC ZrB2z20 vol.-%MoSi2  2.7 6.5  1.361027 3.861028  66 and 86  103  *By TGA up to 1400uC.  Fah renho l tz and H i lmas  Ox ida t ion o f UHTCs  I n t e r n a t i o n a l M a t e r i a l s R e v i ew s  2 0 1 2  VO L 5 7  NO 1  6 9  \\x0c', 'diboride  ceramics.  Both  ZrB2  and  HfB2  undergo  stoichiometric oxidation.  In the  absence of  additives,  the oxide scale that y1100uC is protective kinetics. Modelling supports the conclusion that oxygen  forms  at  temperatures  below  leading to parabolic mass gain  diffusion through the  liquid B2O3  is  the  rate  limiting  step. At  higher  temperatures,  the  B2O3  evaporates,  leaving  a non-protective porous ZrO2 or HfO2  scale  and  leading  to  rapid,  linear mass  gain  kinetics. The  addition of SiC is  the most  common method used to  improve the oxidation behaviour of  the diborides. The  presence  of  SiC extends  the  temperature  range  over  which a protective glassy oxide is present by promoting  the  formation of a silica containing amorphous oxide  layer. The resulting scale has a complex layered structure  that consists of: an outer borosilicate glassy layer, which  provides  oxidation  protection;  a  layer  that  contains  crystalline ZrO2 or HfO2 plus amorphous borosilicate  glass; and a partially oxidised layer that contains porous  crystalline ZrO2 or HfO2 and/or a SiC depleted diboride  layer. The outer borosilicate glassy temperatures of 1600uC or higher. revealed the formation of bubbles  layer  is  stable  to  In situ observations  followed  by  scale  rupture,  confirming  thermodynamic  predictions  that  oxidation protection would be lost at elevated tempera tures due to gas pressure behind the scale. Arc heater  testing has  shown that  the ZrO2  layer can densify and  become protective at extreme conditions  that  simulate  hypersonic flight or atmospheric re-entry. The oxidation  protection of diborides  can also be  improved by  the  addition  of  other  silica  scale  forming  compounds  including MoSi2, Si3N4, and Ta5Si3. Although none of  these compounds alone provides protection equivalent  to the addition of SiC,  the protection can be improved  by compounds such as other transition metal diborides  that form solid solutions with the ZrB2 or HfB2 matrix,  which can lead to changes  in the morphology of  the  oxide scale such as phase separation of  the borosilicate  glass. Finally,  the  oxidation  protection  of  ultra-high  temperature diborides was improved without the use of  silica scale formers such as SiC or MoSi2 through the use  of  additives  that  affected  the  morphology  of  the  crystalline oxide scale.  Reviewing published research on diboride oxidation  has provided a perspective on needs for future research.  As noted at several points in the text, past research has  focused  largely  on  the  composition,  thickness,  and  thermal  stability of  the outer borosilicate glassy/liquid  layer.  The  addition  of  SiC  particles  to  ultra-high  temperature diborides produces a borosilicate layer that least 1600uC. The glassy oxide layer acts as an oxygen diffusion barrier and provides passive  is  stable up to at  oxidation with parabolic mass gain kinetics. However,  the  formation of an outer glassy layer also has  some  negative  impacts  on  performance  of  diboride  based  UHTCs. For  example,  some authors have noted SiC  depletion beneath the outer glassy layer, which produces  significant porosity in the underlying ceramic and may  have an adverse effect on the performance of the ceramic  in extreme  environments.  In addition,  the presence of  SiC leads to the development of residual stresses in the  ceramic  due  to  the mismatch  in  thermal  expansion  coefficients between SiC and the matrix (ZrB2 or HfB2).  These stresses can have a negative impact on the thermal  cycling behaviour and thermal stability of the diborides.  Manipulating  the microstructure  and,  possibly,  the  composition of  the crystalline oxide layer holds promise  for  improving  the  oxidation  behaviour  of  diboride  UHTCs. The use of additives  such as W and transition  metal diborides that  form solid solutions with the ultra high temperature diborides improves oxidation protection  of ultra-high temperature diborides without  forming an  outer SiO2 layer. Further research is needed to understand  the mechanism(s)  by  which  these  additions  improve  oxidation behaviour, which could lead to optimisation of  oxidation protection using this approach. Additives that  alter the structure of  the crystalline oxide layer may also  provide  improved  oxidation  behaviour  in  the  extreme  environments associated with hypersonic flight and atmo spheric re-entry since ZrO2 and HfO2 are more stable than  the glassy oxides that provide protection for SiC contain ing materials at intermediate temperatures. The modifica tion of the crystalline oxide scale is a promising approach  that could produce further gains in oxidation protection.  Another possible avenue of future research is the develop ment of engineered surface structures to enhance oxidation  protection. All of  the  reports discussed in this  review  allowed the surface oxide to develop naturally. In contrast,  improved  oxidation  protection  may  be  possible  by  designing a layered surface structure that would provide  oxidation protection while minimising adverse effects due  to chemical or mechanical  incompatibility. So, although  the oxidation behaviour of ultra-high temperature dibor ides has been studied for many years, further research may  lead to greater understanding of oxidation mechanisms,  better  control  of microstructure  development  in  the  layered oxide  scale, and improved oxidation behaviour  for this class of materials.  Acknowledgements  Research on UHTCs at Missouri University of Science  and Technology has been supported by a number of  grants  over  the  years. Currently, UHTC research  is  supported through the US Air Force Office of Scientific  Research/NASA National Hypersonic  Science Center  for Materials and Structures  (grant no. FA9550-09-1 0477),  the US National  Science  Foundation  (grant  no. DMR-0906584),  and the US Air Force Office of  Scientific Research (grant no. FA9550-09-1-0168). The  authors wish to thank graduate  student Ms Maryam  Kazemzadeh Dehdashti, who provided the SEM images  in Fig. 2. In addition,  the authors are thankful  for the  contributions  of  Shi C. Zhang  and  former  graduate  students Alireza Rezaie and Adam Chamberlain.  References  1. M. J. Gasch, D. T. Ellerby and S. M. Johnson:  in ‘Handbook of  ceramic composites’,  (ed. N. P. Bansal), 197-224; 2005, Boston,  MA, Kluwer Academic Publishers.  2. A. L Chamberlain, W. G. Fahrenholtz, G. E. 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},{
  "_id": 173,
  "PDF": "Oxidation of zirconium diboride with niobium additions.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 33 (2013) 1591-1598  Oxidation of zirconium diboride with niobium additions  ∗  Maryam Kazemzadeh Dehdashti  , William G. Fahrenholtz, Greg E. Hilmas  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65401, United States  Received 20 January 2013; accepted 26 January 2013  Available online 17 February 2013  Abstract     Oxidation of ZrB2 ceramics containing Nb additions at 1500 C resulted in the formation of a two-layer oxide scale. The outer surface was partially covered by a glassy layer containing B2O3 with smaller amounts of Nb and Zr oxides dissolved into it. With increasing exposure time, evaporation of B2O3 from the outer layer resulted in precipitation of oxide particles in the receding glassy phase. Between the outer layer and the unoxidized (Zr,Nb)B2 was a porous layer that consisted of particles containing Zr, Nb, and O. The formation of Nb2Zr6O17 was observed in the porous oxide layer. Since this compound is solid at the oxidation temperature, liquid phase sintering of the ZrO2 scale was not possible. However, dissolution of Nb into B2O3 increased the stability of the liquid/glassy layer, which acted as a barrier to the transport of oxygen at higher temperatures compared to the scale formed on nominally pure ZrB2 . © 2013 Elsevier Ltd. All rights reserved.  Keywords: ZrB2 ; Nb; Oxidation; ZrO2 ; Composites  1.   Introduction  Ultra high  temperature ceramics  (UHTCs) are a group of materials that includes ZrB2 , ZrC, HfB2 and HfC. These materials are candidates  for applications  that  require exposure  to extreme thermal and chemical environments. The performance advantages of  the diboride-based UHTCs come not only from their high-temperature stability but also from  the capability  to transfer and redistribute heat at elevated temperatures. This characteristic  is attractive  for  sharp  leading edges  for hypersonic aerospace vehicles, which must  transfer heat away  from  the hottest areas and redistribute it to cooler areas.1 Among UHTCs, ZrB2 has the lowest theoretical density combined with reported 100 W/m K at  thermal conductivity values as high as  room temperature, which  is an advantage over other candidates  for aerospace applications.2 Oxidation behavior  is a  restriction  to  the development of ZrB2 -based ceramics for rocket propulsion and hypersonic flight applications. Assuming stoichiometric oxidation according  to Reaction (1), exposure of ZrB2 to air at temperatures of 800 C     ∗  Corresponding author. Tel.: +1 573 578 2853; fax: +1 573 341 6934.  E-mail addresses: mk7y8@mst.edu (M. Kazemzadeh Dehdashti),  billf@mst.edu (W.G. Fahrenholtz), ghilmas@mst.edu (G.E. Hilmas).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.01.033  and above results in formation of B2O3 and ZrO2 , which leads to measurable mass gain.3  ZrB2(cr) +   5/2 O2(g) →   ZrO2(cr) +   B2O3(l)  (1)        Evaporation of B2O3 is considerable at  temperatures above 1200 C. The  loss of B2O3(l) leaves behind a porous ZrO2 layer with a columnar microstructure, which offers channels for rapid oxygen  transport  to  the reaction  interface and results in significant mass gain at  temperatures above 1200 C.4 The conventional approach to improving the oxidation resistance of diboride ceramics  is  to add Si-containing compounds such as SiC1,5-13 , MoSi2 14-16 , or TaSi2 .17,18 The formation of a borosilicate  layer on  the surface of  the diborides provides  improved oxidation resistance in air compared to the borate glass on nominally pure ZrB2 due to the increased stability of the borosilicate glass compared to the borate material.3,19 At elevated temperatures, SiO(g) forms beneath  the borosilicate glass as a result of active oxidation of SiC. When  the pressure of SiO(g) exceeds ambient, the resulting pressure can rupture the protective glassy layer, which  can  result  in  a  cyclic protective/non-protective scale-forming sequence.20 Further, some authors have noted the formation of the SiC-depleted layer in ZrB2 -SiC samples, which facilitates the transport of oxygen through the oxide scale to the unoxidized matrix.21,22 Hence, SiC may not be the best choice        \\x0c', '1592   M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598  for improving the oxidation resistance of ZrB2 ceramics at ultra high temperatures. Several studies showed  that additions of Cr-, Ti-, Nb-, V-, and Ta-borides improved the oxidation resistance of ZrB2 -SiC composites.22-26 Hence,  additions of  transition metals offer improved oxidation  resistance  to ZrB2 ceramics without  the deleterious  effects of  silica  formers. Similarly, Zhang  et  al. reported  that WC additions  improved  the oxidation  resistance of ZrB2 ceramics. The addition of WC  led  to  the formation of a  two-layer scale structure, which consisted of a porous zirconia outer  layer and a dense  inner  layer containing ZrO2 and WO3 ,  in contrast with  the single, highly porous and columnar ZrO2 layer  formed on nominally pure ZrB2 .27,28 It was suggested that during oxidation, the presence of WO3 in the oxide scale resulted in liquid phase sintering of ZrO2 , which increased the relative density of the scale, resulting in improved oxidation resistance.28 Other  transition metal additives, such as Nb, may offer similar beneficial effects on the oxidation behavior of zirconium diboride. Unlike the presence of WO3 , the formation of Nb2O5 during oxidation  is not expected  to result  in  the formation of a  liquid phase. Examination of  the Nb2O5-ZrO2 phase diagram shows that at 1500 C, the presence of small (less than 10 mol%) concentrations of Nb should lead to the formation of solid compounds such as Nb-doped ZrO2 or Nb2Zr6O17 rather than a liquid phase, such as the WO3-ZrO2 solution predicted for the presence of small amounts of W with ZrO2 .29,30 The purpose of this paper was to examine the effect of niobium additions on the oxidation behavior of ZrB2 at elevated temperatures to gain insight into the mechanism by which transition metal additions improve oxidation resistance.     2. Experimental procedure  High purity  (>99%) ZrB2 powder  (Grade B, H.C. Starck, 2  Newton, MA) with an average particle size of  \\u242em was used to prepare  the specimens for  this study. To enhance densification, 2 wt% B4C (0.8  \\u242em, Grade HS, H.C. Starck) was added to all batches  to react with and remove oxide  impurities from the powder particle surfaces. For some batches, 6 mol% niobium was added in form of Nb powder (Johnson Matthey, MA), 1  which had an average particle size of  \\u242em. To reduce particle size and promote  intimate mixing,  the as-received ZrB2 , B4C, and Nb (for Nb containing batches) were dispersed  in methyl ethyl ketone  (MEK) by ball milling with zirconia media  for 24 h. An organic dispersant (DISPERBYK-110, BYK-Chemie Co., Wesel, Germany) was added at a  level of 0.54 mg of dispersant per m2 of ZrB2 surface area. The amount of zirconia contamination added  to  the batches as a  result of ball milling was determined to be less than 1 wt% based on the mass of ZrB2 powder by measuring  the mass of  the media before and after milling. After mixing,  the  solvent was  removed using  rotary −80 evaporation, and then the powder was ground and sieved to  mesh. Powders were densified by hot pressing  (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 2100 C for 45 min at a pressure of 32 MPa. Powders were loaded into a graphite die lined with graphite foil  that was coated with BN spray. Billets     25 mm and a thickness of  5 mm were prowith a diameter of  duced. Specimens with dimensions of 10 mm by 4 mm by 4 mm were diced from the billets and polished on all sides to a 15  \\u242em finish for testing and characterization. Images obtained by scanning electron microscopy (SEM; S-4700, Hitachi, Japan) from the polished surfaces of as processed (Zr,Nb)B2 and ZrB2 were used to study the microstructure of the specimens. The amount of B4C remaining after densification was calculated using  image analysis software  (ImageJ, U.S. National  Institutes of Health, Bethesda, MD). The bulk densities of  the hot pressed billets were measured using  the Archimedes  technique with water as the immersing medium. Oxidation  studies were performed  in a MoSi2 resistanceheated  horizontal  tube  furnace  (Model  0000543  Rapid Temperature Furnace, CM Inc., Bloomfield, NJ) equipped with a high-purity alumina tube with a diameter of 6.35 cm. Specimens were cleaned  in acetone  in an ultrasonic bath and  then placed on a zirconia foam setter  that was on an alumina D-tube. The specimen assembly was  inserted  into  the center of  the furnace and  leveled. The ends of  the  tube were sealed using gas-tight end caps. Specimens were heated at  C/min  to 1500 C or 1600 C and held for up to 3 h in air with a flow rate of 0.2 cm/s (linear flow rate was calculated according to the volumetric flow rate and the size of the tube). To minimize changes such as further oxidation  that may occur during cooling, specimens were air quenched  to room  temperature by removing  them from  the furnace after the desired oxidation time. The  thicknesses of  the resulting oxidation  layers were measured  from  fracture  surfaces  that were  observed  in  SEM. In  addition,  the microstructures  of  the  oxide  scales were observed using SEM and chemical compositions were analyzed using energy dispersive spectroscopy (EDS; EDAX, Mahwah, NJ). X-ray diffraction  (XRD; Philips X-Pert Pro diffractometer, Westborough, MA)  analysis was used  to  identify major crystalline phases present in both the pre-oxidized and the postoxidized composites and  the data were analyzed using X‘Pert High Score software.  5           3. Results and discussion  3.1. Densification, microstructure, and phase analysis  A microstructure typical of the (Zr,Nb)B2 specimens used in this investigation is presented in Fig. 1. The darker phase is B4C and it appears to be uniformly dispersed in the lighter (Zr,Nb)B2 matrix. Based on image analysis, the amount of B4C remaining after densification was 1.4 wt%. The average bulk density of bars cut from hot-pressed (Zr,Nb)B2 billets was 6.03 g/cm3 . Using a volumetric rule of mixtures calculation, and assuming true densities of 6.09 g/cm3 for ZrB2 , 2.52 g/cm3 for B4C, and 8.57 g/cm3 for Nb,  the  theoretical density of ZrB2 containing 6 mol% Nb to be 6.05 g/cm3 . Using  was calculated  this  true density,  the hot-pressed bars had  relative densities of >99%. The calculated relative density  is consistent with  the minimal amount of porosity revealed by SEM analysis. Also, Archimedes’ measurements showed  the amount of open porosity  to be  insignificant. Thus, porosity was not considered to have a significant effect on  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598   1593  Fig. 1. SEM   image of a polished cross section showing   the microstructure of  ZrB2 containing 1.5 vol% B4C and 6 mol% Nb.  the oxidation behavior. Microstructure and phase analysis using SEM and XRD were consistent with the dissolution of Nb into the matrix  to form a (Zr,Nb)B2 solution, which was  identified as hexagonal by indexing to ZrB2 (PDF card number 34-0423).  3.2. Surface morphology and composition        Assuming that the oxidation of (Zr1−xNbx )B2 proceeds stoichiometrically,  reaction at  temperatures of 800 C and above 450 should produce molten B2O3 (melting  temperature  C), solid ZrO2 , and solid Nb2O5 in the molar ratios shown in Reaction  (2).  In  this case,  the addition of 6 mol% Nb  to ZrB2 is equivalent  to x = 0.06, which produces an oxide  scale with a molar  ratio of Nb2O5 to ZrO2 of 1  to 33 or 93.5 wt% ZrO2 plus 6.5 wt% Nb2O5 . According  to  the ZrO2-Nb2O5 phase diagram29,30 ,  the primary crystalline phases  that should  form are a monoclinic  solid  solution based on ZrO2 that contains dissolved Nb and an orthorhombic compound, Nb2Zr6O17 . By increasing  the  temperature,  the solubility  limit of Nb2O5 into ZrO2 increases and  the  ratio of Nb2Zr6O17 to  the ZrO2 solid solution decreases. The melting  temperature of Nb2Zr6O17 is 1670 C and  the solubility  limit of Nb2O5 in ZrO2 at 1500 C is about 7 mol% (14 mol% Nb). Due  to  the formation of solid Nb2Zr6O17 , no  liquid phase  is predicted  for  the composition of  the oxide scale, which should be ZrO2 containing 6 mol% of dissolved Nb at 1500 C. Upon cooling  to  room  temperature, the solubility limit of Nb2O5 in ZrO2 decreases and some Nb2Zr6O17 should precipitate  from  the ZrO2 solid  solution. At  room  temperature, Nb2Zr6O17 comprises  about 18 mol% of  the oxide phase  along with  an  amorphous phase  that  is mainly B2O3 . (Zr1−xNbx )B2(cr) +  (1−x) ZrO2(cr)  x/4) O2(g) →   (5/2   +           +   x/2 Nb2O5(cr) +   B2O3(l)  (2)  The surface of  (Zr,Nb)B2 oxidized at 1500 C  for 0 h  (i.e., quenched as soon as it reached 1500 C) is shown in Fig. 2a. A majority of  the surface of  the specimen was covered by a dark phase that had a glassy appearance (90% of the surface area)        Fig. 2. Surface of (Zr,Nb)B2 oxidized at 1500 the presence of a glassy oxide (dark phase) and crystalline oxide particles (light  C for (a) 0 h, (b) 3 h, showing     phase).     with a small fraction of the surface composed of oxide particles (10% of  the surface area). After 3 h at 1500 C (Fig. 2b),  the 60% and area  fraction of  the glassy phase had decreased  to  the glassy phase was concentrated  in several pools  that were surrounded by oxide particles. Due  to  low sensitivity of EDS  to  light elements, quantification of  the boron content  in  the glassy phase was not possible. However, EDS  results  indicated  that  the matrix of  the glassy phase contained O and Nb along with a small amount of Zr, presumably all dissolved  in B2O3 . Small particles  that contained both Zr and Nb (Fig. 3a) were also observed in the glassy phase. 31 and Nb2O5-B2O3 32 phase diaAccording to the ZrO2 -B2O3 grams, approximately 12 mol% ZrO2 can dissolve into B2O3 at 1500 C while both Nb2O5 and B2O3 are  liquids at  that  temperature. Hence,  the particles observed at  room  temperature could be ZrO2 , Nb2O5 , or an oxide containing Zr and Nb. The particles could have  formed either during oxidation or during cooling. They could form during oxidation due  to evaporation of B2O3 that would result  in supersaturation of  the remaining B2O3 with ZrO2 , which could result in precipitation. Conversely,         \\x0c', '1594   M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598  Fig. 3. SEM images of the surface of a (Zr,Nb)B2 specimen oxidized at 1500 for 3 h that were (a) close to the edge of the liquid pool, and (b) in the middle of  C     Fig. 4. SEM images of a (Zr,Nb)B2 specimen oxidized at 1500 the edge of the liquid pool, (b) in the crystalline oxide region close to the liquid  C for 3 h (a) at     pool, showing the growth of the oxide particles by joining the small precipitated  the liquid pool. The images show a homogeneous distribution of oxide particles  particles.  containing Zr and Nb and formation of elongated particles.  (Zr1−xNbx )O2+0.5x  the particles could precipitate from  the glassy phase when  the specimen was cooled  from  the processing  temperature due  to the change  in  solubility of ZrO2 in B2O3 with  temperature. According  to  the ZrO2-Nb2O5 phase diagram29 ,  two different crystalline phases could form when ZrO2 and Nb2O5 precipitate from the B2O3 melt. For Nb2O5 contents less than about 5 mol%, a  solid  solution  is  the  stable phase.  If  the Nb2O5 concentration in the glassy phase is higher, then the crystalline phase Nb2Zr6O17 (also designated 6ZrO2 ·Nb2O5 ) could form in addition to the ZrO2 solid solution. X-ray diffraction was used to characterize the phases present on the surface of oxidized (Zr,Nb)B2 . The major phases detected were  triclinic H3BO3 , monoclinic ZrO2 and orthorhombic Nb2Zr6O17 (PDF card numbers 30-0199, 83-0944 and 72-1745, respectively). Boric acid formed after cooling  the specimen  to room  temperature due  to the instability of B2O3 in the humid ambient air. With increasing exposure time, more B2O3 should evaporate from the surface, which would increase precipitation of Zrand  Nb-rich oxide particles  from  the  liquid phase. Near  the edges of the pools, spherical particles precipitated as the B2O3 liquid evaporated. Fig. 3a shows that the particles were uniformly distributed in the glassy phase and that they had an average diameter \\u242em. Closer to the centers of the glassy pools, elonof about 0.5  gated precipitates appeared to grow from the spherical particles (Fig. 3b). The elongated precipitates were typically about 3  \\u242em long. As  the glassy phase  receded during extended exposure at 1500 C,  the underlying oxide particles were revealed (Fig. 4a and b). As the liquid receded, the small spherical particles that were observed in the glassy phase appeared to attach to the larger particles  that were exposed. Some smaller particles were visible between  the  larger ones.  In areas where  the glassy phase remained, particularly in the middle of the glassy pools, the precipitated particles became more concentrated. As can be seen in Fig. 5, clusters of elongated particles formed in the areas where the last of the glassy pools were present.     \\x0c', 'M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598   1595  Fig. 5. SEM   (Zr,Nb)B2 showing the clustering of the equiaxed and elongated particles on the surface of  specimen oxidized at 1500  image of   for 3 h  the   C      the crystalline oxide particles.  3.3. Oxide scale morphology and thickness        Initial attempts to polish cross sections of oxidized samples revealed that the oxide scales were damaged by the preparation process. To produce cross sections  that were representative of the oxide scales, fracture surfaces were examined. Fig. 6 shows a low magnification SEM image of the fracture surface of a (Zr,Nb)B2 sample oxidized at 1500 C for 3 h. Significant differences were observed in the thickness of the glassy layer between the middle and edge of the glassy pool. After oxidation at 1500 C for 3 h the thickness of the glassy layer ranged 40  \\u242em.  from a maximum of about 0  to  In addition,  the \\u242em thickness of  the porous oxide  layer ranged from about 45  to 60  \\u242em. Areas of thick glassy oxide had thinner porous oxide layers while areas with thinner glassy oxide had thicker porous layers. Fig. 7 shows cross sectional SEM images of the oxide scales from regions with maximum glassy layer thickness. The scales were  formed on  the  surfaces of  the  (Zr,Nb)B2 samples after oxidation at 1500 C for 0, 1.5 and 3 h. The oxide scales in these regions consisted of  two  layers: (1) a dense outer glassy  layer, and (2) an inner layer that appeared to be porous. The two-layer scale  is believed  to be  formed due  to volume expansion upon 300% volume conversion of ZrB2 to ZrO2 and B2O3 , which is   \\u242em      Fig. 6. Fracture surface of a (Zr,Nb)B2 specimen oxidized at 1500     C for 3 h.  Fig. 7. Fracture surfaces of (Zr,Nb)B2 oxidized at 1500 and (c) 3 h.     C for: (a) 0 h, (b) 1.5 h,  expansion based on density calculations. Oxidation produces two phases because of  the  immiscibility of  the  two materials while the large volume expansion associated with formation of B2O3 causes it to be forced to the surface of the specimen.3 Several  small  elongated particles  can be observed  in  the glassy layer in Fig. 7a and b. The particles were uniformly distributed  through  the  thickness of  the glassy  layer. Also, some residual glassy phase was observed between the oxide particles in  the porous  layer. The oxide particles  in  the porous  layer of oxidized (Zr,Nb)B2 were  less  than 10  in diameter and had equiaxed shapes. For comparison,  the scale  formed on nominally pure ZrB2 was composed of larger ZrO2 particles that had an elongated morphology.33 Fig. 8 shows the results of the scale thickness measurements for nominally pure ZrB2 compared  to  (Zr,Nb)B2 after oxidation at 1500 C for 0, 1.5 and 3 h. While  the scales formed on nominally pure ZrB2 were mostly uniform, considerable differences  in  the  thickness of  the scales between  the middle and edge of  the glassy pools were observed  in  the (Zr,Nb)B2 specimens. Therefore,  the measurements were performed on  the regions with maximum glassy  layer  thickness  for  (Zr,Nb)B2 . When nominally pure ZrB2 and  (Zr,Nb)B2 reached 1500 C,  \\u242em         \\x0c', '1596   M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598     \\u242em   increased the stability of the liquid phase at high temperatures. Dissolution of Nb2O5 into the liquid phase should decrease the activity of B2O3 and, consequently, reduce its vapor pressure and evaporation rate. The presence of a glassy layer should improve the oxidation resistance of the ceramic since borate glasses act as a barrier to oxygen transport.34 After  oxidation  75  at  1500 C  for  3 h,  the  thickness  of  the porous  layer was  \\u242em for nominally pure ZrB2 compared 45  to  for  (Zr,Nb)B2 . Based on  this observation,  the stability of  the glassy  layer resulted  in a decreased oxidation rate for (Zr,Nb)B2 . The addition of Nb increased the stability of the liquid phase, providing better protection at high  temperatures, which decreased the oxidation rate of the (Zr,Nb)B2 compared with nominally pure ZrB2 . Not only was the initial thickness of the porous oxide layer thinner for (Zr,Nb)B2 (18  26  \\u242em compared to  \\u242em for nominally pure ZrB2 ), the porous oxide layer was C (45  75  thinner after 3 h at 1500 \\u242em compared to  \\u242em for nominally pure ZrB2 ). Based on the calculations performed on the thickness of the porous layers, the oxidation exhibited linear kinetic behavior (R2 > 0.94) at 1500 C. Further, the addition of Nb to ZrB2 resulted in a lower oxidation rate (9  \\u242em/h) compared to nominally pure ZrB2 (16  \\u242em/h).        3.4. Evolution of structure  Fig. 9 is a schematic description of the evolution of the structure of the oxide scale on (Zr,Nb)B2 . At the early stages, Zr, Nb, and B oxides form as (Zr,Nb)B2 is oxidized. The Zr-Nb oxide particles form on the (Zr,Nb)B2 surface, but are covered by a liquid phase composed of mainly B2O3 with smaller amounts of dissolved Nb and Zr oxides. During oxidation, Nb and Zr oxides dissolve  into  the  liquid borate phase. At higher  temperatures, the evaporation rate of B2O3 from  the outer surface  increases, which results in precipitation of particles in the liquid. The particles contain Zr, Nb, and O and join together to form either the equiaxed or elongated grains visible  in  the glassy phase of  the quenched specimens (e.g., Fig. 3). As exposure time increases, B2O3 evaporation continues and  the  liquid phase concentrates in pools that are separated by regions of crystalline oxide. As the glass recedes, it exposes the underlying porous oxide scale. Previous studies have shown that the addition of W to ZrB2 leads to liquid phase sintering of the porous ZrO2 scale, which improved the oxidation resistance by decreasing oxygen transport through the scale. In contrast, Nb2O5 and ZrO2 form a solid Zr-Nb-O  Fig. 8. Scale   thickness as a function of oxidation   time at 1500     C comparing  nominally pure ZrB2 to areas with the maximum glassy thickness for (Zr,Nb)B2 .     \\u242em   4  the glassy  layers  for both materials were  thick, while 26  18  the porous oxide scale was  \\u242em for ZrB2 and  \\u242em for (Zr,Nb)B2 . After 1.5 h at 1500 C, the glassy and porous layers 2 and  59  on the nominally pure ZrB2 were  \\u242em thick, respectively. For the same oxidation time, (Zr,Nb)B2 had a maximum 18  glassy  layer  thickness of  \\u242em. The porous  layer beneath 35  the glassy layer with the maximum thickness was  \\u242em. No glassy  layer was observed on pure ZrB2 after 3 h at 1500 C, while the glassy layer on (Zr,Nb)B2 had a maximum thickness 37  of  \\u242em. The thickness of the porous layer on nominally pure 73  45  ZrB2 was  \\u242em compared  to  for  (Zr,Nb)B2 after 3 h at 1500 C. Although  the surface area of  the glassy pools on  (Zr,Nb)B2 decreased with  increasing  the exposure  time at 1500 C, the thickness of the glassy layer increased with exposure time. For nominally pure ZrB2 ,  the  thickness of  the glassy  layer decreased with  increasing oxidation  time  at 1500 C due  to evaporation of B2O3 . Previous studies have reported that evaporation of B2O3 is substantial above 1100 C due  to  its high vapor pressure.3 In contrast,  increasing  the oxidation  time at 1500 C for (Zr,Nb)B2 resulted  in an  increase  in  the  thickness of the glassy layer. The presence of glassy phase after oxidation of (Zr,Nb)B2 at 1500 C for 3 h indicated that the addition of Nb  \\u242em                        Fig. 9. Model of the evolution of the oxide structure of (Zr,Nb)B2 .  \\x0c', 'M. Kazemzadeh Dehdashti et al. / Journal of the European Ceramic Society 33 (2013) 1591-1598   1597  compound at the oxidation temperature, namely Nb2Zr6O17 , so no  liquid phase sintering of ZrO2 occurred. However, dissolution of Nb2O5 into the B2O3 liquid phase increased the stability of  the  liquid phase compared  to  the nominally pure B2O3 that forms when ZrB2 is oxidized. The  improved  stability of  the glassy  layer  leads  to  improved oxidation behavior because  the external layer and glassy phase trapped among the particles that make up the porous oxide layer acts as a barrier to the transport of oxygen. Therefore, the addition of Nb to ZrB2 increased the oxidation resistance, but only as long as the Nb-containing B2O3 phase was present.  4. Conclusion     The oxidation behavior of (Zr,Nb)B2 ceramics was studied. At 1500 C, exposure to air resulted in the formation of a twolayer oxide scale structure on (Zr,Nb)B2 . The two layers were: (1) an outer  layer of a glassy phase containing B2O3 with Nb and Zr dissolved  in  it; and (2) a porous oxide  layer composed of oxide particles containing Zr and Nb. Small spherical particles, presumably a ZrO2 containing dissolved Nb, grew  in  the glassy phase with increasing exposure time. Some of the particles were spherical while others were elongated. As  the B2O3 evaporated,  the particles became concentrated and were eventually  incorporated  into  the newly exposed porous oxide  layer, which contained both ZrO2 and Nb2Zr6O7 . As the glass receded, the small precipitated particles  joined  the porous oxide  layer that was present under  the glassy  layer. Because of high melting point of Nb2Zr6O7 that was  formed  in  the porous oxide layer, liquid phase sintering was not active as has been reported for W-containing ZrB2 . However, dissolution of Nb  into  the B2O3 liquid phase  increased  the stability of  the protective  liquid  layer by  reducing  the volatility of B2O3 from  the  liquid phase. Hence (Zr,Nb)B2 showed improved oxidation resistance compared with pure ZrB2 .  Acknowledgments  This work was supported as part of the National Hypersonic Science Center for Materials and Structures (Grant FA9550-091-0477) with Dr. Ali Sayir (AFOSR) and Dr. Anthony Calomino (NASA) as program managers. The authors wish to thank project principal investigator Dr. David Marshall of Teledyne Scientific and Imaging for his support and guidance.  References  1. Monteverde F, Savino R, Fumo MDS, Di Maso A. Plasma wind   tunnel  testing of ultra-high  temperature ZrB2 -SiC composites under hypersonic re-entry conditions. J Eur Ceram Soc 2010;30(11):2313-21.  2. Guo SQ. Densification of ZrB2 -based composites and their mechanical and physical properties: a review. J Eur Ceram Soc 2009;29(6):995-1011.  3. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during   the  oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27(6):2495-501.  4. Parthasarathy TA, Rapp RA, Opeka M, Kerans RJ. Effects of phase change  and oxygen permeability in oxide scales on oxidation kinetics of ZrB2 and HfB2 . J Am Ceram Soc 2009;92(5):1079-86.     C. J Eur  5. Rezaie A, Fahrenholtz WG, Hilmas GE. 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Peng F, Berta Y, Speyer RF. Effect of SiC, TaB2 and TaSi2 additives on the isothermal oxidation resistance of fully dense zirconium diboride. J Mater Sci 2009;24(5):1855-67.  24. Talmy IG, Zaykoski JA, Opeka MM, Dallek S. Oxidation of ZrB2 ceramics modified with SiC and group IV-VI transition metal diborides. Electrochem Soc Proc 2001;12:144-58.  25. Hu P, Zhang XH, Han JC, Luo XG, Du SY. Effect of various additives on  the oxidation behavior of ZrB2 -based ultra-high-temperature ceramics at 1800 C. J Am Ceram Soc 2010;93(2):345-9.     26. Zhang XH, Hu P, Han JC, Xu L, Meng SH. The addition of   lanthanum  hexaboride  to zirconium diboride  Mater 2007;57(11):1036-9.  for   improved oxidation   resistance. Scr  27. Zhang SC, Hilmas GE, Fahrenholtz WG.   Improved oxidation   resistance  of  zirconium diboride by  2008;91(11):3530-5.  tungsten   carbide   additions.   J Am Ceram Soc  28. Zhang SC, Hilmas GE, Fahrenholtz WG. Oxidation of zirconium diboride with tungsten carbide additions. 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},{
  "_id": 174,
  "PDF": "Oxidation of Zirconium Diboride with Tungsten Carbide Additions.pdf",
  "Text": "['Oxidation of Zirconium Diboride with Tungsten Carbide Additions  w,\\x03 Greg E. Hilmas,\\x03\\x03 and William G. Fahrenholtz\\x03\\x03 Shi C. Zhang,  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, Missouri 65409  The oxidation resistance of zirconium diboride (ZrB2) ceramics with tungsten carbide (WC) additions was studied. ZrB2 ceramics, nominally pure and with WC additions of 4, 5, and 6 mol%, were densified by pressureless sintering. During heating, the WC that was added to the ceramic dissolved into the ZrB2 matrix, forming a solid solution. Oxidation behavior was evaluated using isothermal furnace oxidation at 15001C and 16001C. The presence of tungsten reduced both the weight gain and oxide scale thickness compared with nominally pure ZrB2. After oxidation at 16001C for 3 h, the thickness of the ZrO2 outer scale was reduced from 4.7 mm for nominally pure ZrB2 to 0.5 mm for the ceramic containing 6 mol% WC. Analysis showed that improved the oxidation resistance of ZrB2 by changing the morphology of the zirconia scale formed during oxidation.  WC additions  I.  Introduction  Z IRCONIUM DIBORIDE (ZrB2)1,2 has been proposed for a variety of aerospace applications, such as leading edges and thermal protection systems for reusable atmospheric reentry vehicles vehicles.3-5 As with  and  hypersonic  flight  other  nonoxide  ceramics, ZrB2 oxidizes when exposed to air at elevated temperatures.6,7 The oxidation behavior of ZrB2 has been studied since the 1960s.8-12 In general, the oxidation of ZrB2 can be considered to be stoichiometric according to Reaction 1, resulting in measurable mass gain above about 8001C.13  ZrB2 ðcrÞ þ 5=2 O2 ðgÞ ! ZrO2 ðcrÞ þ B2O3 ðlÞ  (1)  Above B12001C, evaporation of B2O3 hind a porous ZrO2 layer that does not protect the underlying from further oxidation. Recently, Parthasarathy et al.14 ZrB2 proposed a model to describe the oxidation behavior and final  is rapid,  leaving be oxidized microstructures for ZrB2, TiB2, and HfB2. The model, and comparison with experimental results,5,8,12,15 revealed that  the ZrO2 scale was not protective at elevated temperatures because it was porous. The porosity, combined with the columnar  microstructure of the ZrO2 scale, offers channels for oxygen transport directly and rapidly to the reaction interface.16 Oxygen  transport is much faster through the resulting pores than diffu sion through the ZrO2 crystal structure would be. Based on the porous scale, the mass gain for diboride ceramics increases rapidly as the temperature approaches 15001C. To improve the oxidation resistance of ZrB2-based ceramics above 14001C, a common approach is the addition of Si-containing compounds,  such as SiC, MoSi2, or TaSi2 into ZrB2 to form an adherent, protective layer of borosilicate glass.17-28 An  alternate  ap proach,  increasing  the density of  the ZrO2  scale, may  also  improve  the oxidation resistance of ZrB2  in this  temperature  range.  In our previous study, ZrB2 containing 4 mol% WC exhibited improved oxidation resistance, compared with nominally pure ZrB2 as measured by weight gain and oxide scale thickness.29 The improved oxidation resistance was attributed to liquid  phase formation between the condensed oxide reaction prod ucts, WO3 and ZrO2, which enhanced the densification of ZrO2 and resulted in the formation of a dense outer scale.  During oxidation of ZrB2 with tungsten carbide (WC) additions, several processes occur simultaneously, including oxidation of WC and ZrB2, liquid phase formation30 between WO3 and ZrO2, densification of ZrO2, and evaporation of B2O3 and WO3. Both WO3 formation and WO3 evaporation are important factors for the persistence of the dense ZrO2 outer scale. If, like B2O3, the WO3 that formed during oxidation evaporated rapidly, then it may not be an effective liquid phase former for sintering  ZrO2. Based on thermodynamic calculations and observations of sintered mixtures of ZrO2 and 4 mol% WO3 reported in the previous study,29 the W-rich oxide phase did not  appear  to  undergo significant evaporation despite its relatively high vapor  pressure. This may be due to the low W concentration in the  system, which would reduce the activity of W in ZrB2 and the activity of WO3 in the resulting oxide liquid. Thus, the actual vapor pressure of the (WO3)n species may be lower than the values predicted using standard state thermodynamic calculations.  This research focuses on optimizing the amount of WC added  to ZrB2 temperature.  ceramics  to improve oxidation resistance at elevated  II.  Experimental Procedures  (1)  Powder Preparation and Characterization  As-received ZrB2 used to prepare the materials for this study. Tungsten was added  (Grade B, H.C. Starck, Newton, MA) was  as WC powder (Cerac Inc., Milwaukee, WI)  to some batches.  The average particle sizes of  the raw powders were calculated  from surface area values determined from nitrogen adsorption  (NOVA 1000, Quantachrome, Boynton Beach, FL) analysis.  The purity, particle size, and suppliers of raw ZrB2, WC powders, are summarized in Table I.  (2)  Specimen Preparation  To enhance densification of ZrB2, 2 wt% B4C (Grade HS, H.C. Starck) was added to all batches to aid in the removal of the surfaces.31 The  oxide  impurities  from the  powder  particle  Table I.  Raw Material Characteristics  Material  Grade  Purity  (wt%)  Particle  Size (mm)  Surface  Area (m2/g)  Oxygen  content  (wt%)  Supplier  ZrB2 B4C WC  B  499  2  1  0.9  H.C. Starck  HS  0.8  15.8  1.3  H.C. Starck  —  499.5  o1  —  —  ACER  ZrB2, zirconium diboride; WC, tungsten carbide.  T. Parthasarthy—contributing editor  This work was  funded by the National Hypersonic Science Center  for Hypersonic  Materials and Structures through a sub-contract from Teledyne Scientific to Missouri S&T. \\x03Member, The American Ceramic Society. \\x03\\x03Fellow, The American Ceramic Society. Author to whom correspondence should be addressed. e-mail: scz@mst.edu  w  Manuscript No. 27876. Received April 18, 2010; approved September 29, 2010.  Journal  J. Am. Ceram. Soc., 94 [4] 1198 - 1205 (2011)  DOI: 10.1111/j.1551-2916.2010.04216.x  r 2010 The American Ceramic Society  1198  \\x0c', 'April 2011  Oxidation of Zirconium Diboride with Tungsten Carbide Additions  1199  Fig. 1. X-Ray diffraction patterns of zirconium diboride (ZrB2) containing 4, 8, and 12 mol% tungsten carbide (WC) after sintering at 19751C showing the peaks shifted to higher values of 2y compared with the powder diffraction file data (vertical lines).  as-received ZrB2, 2 wt% B4C, batches) were dispersed in methyl ethyl ketone by ball milling  and WC (for W-containing  for 24 h with a dispersant  (DISPERBYK-110, BYK-Chemie  Co., Wesel, Germany). Next, 1 wt% binder (QPAC-40, Empower  Materials, Newark, DE) was added and the mixture was milled  for another 24 h. After ball milling, the slurry was dried, ground  and sieved to 50 mesh for dry pressing. Cylindrical disks of 19 mm  in diameter were  formed by uniaxial pressing  at  18.6 MPa  (2.7 ksi) followed by cold isostatic pressing at 315 MPa (45 ksi).  Compacted pellets were densified by pressureless sintering. Pellets were heated at 101C/min in a graphite crucible to 19751C  using a graphite element  furnace (3060-FP20, Thermal Tech nology, Santa Rosa, CA). The furnace atmosphere was a mild vacuum (B20 Pa) for temperatures of 16501C or below, but was switched to flowing argon (B105 Pa) for temperatures above 16501C. Hold times ranged from 2 to 3 h for  sintering. The  phases present after heating were analyzed by X-ray diffraction  (XRD; XDS 2000, Scintag Inc., Cupertino, CA). The bulk den sities of sintered specimens were determined using Archimedes’  method with water as the immersing medium. Relative densities  were calculated by dividing the measured bulk densities by the  B5 mm were produced. Specimens with dimensions of 5 mm by 4 mm by 3 mm were diced from the pellets for oxidation tests.  (3)  Oxidation Testing  The weight gain was characterized as a function of 15001C or 16001C using static  testing in a MoSi2 heated horizontal mullite tube furnace (Model 0000543 Rapid  time at  resistance  Temperature Furnace, CM Inc., Bloomfield, NJ). The tube had  an inside diameter of 7.5 cm and was fitted with gas-tight end  caps. Before oxidation, specimens were cleaned in an ultrasonic  bath with acetone. Specimens were placed on ridged zirconia  setters to minimize the contact area between the specimens and at B51C/min setters. Specimens were heated to 15001C or 16001C and held for 1, 2, or 3 h in air with a flow rate of 1 l/  min. After cooling to room temperature, mass was measured (mg/cm2) were calculated  again. The normalized weight gains  from mass gain and the surface area was calculated from the  specimens based on their measured dimensions.  (4)  Sample Characterization  theoretical densities, which were calculated for each composition  The thicknesses of the resulting oxidation layers were measured  based on the nominal amounts of B4C and WC in the batches. Densified pellets with a diameter of B18 mm and thickness of  from polished cross-sections. The microstructures of  the oxide  scales were observed using scanning electron microscopy (SEM;  Fig. 2. Images of nominally pure zirconium diboride (ZrB2) and ZrB2 ceramics with tungsten carbide (WC) additions of 4, 8, and 16 mol% after isothermal oxidation at 16001C for 180 min.  \\x0c', '1200  Journal of the American Ceramic Society—Zhang et al.  Vol. 94, No. 4  Fig. 3.  Comparison of weight gains for nominally pure zirconium dibo ride (ZrB2) (circles) and ZrB2 ceramics with WC additions of 4 (squares), 5 (diamonds), and 6 (triangles) mol% oxidized at 16001C for various  hold times.  Hitachi S-570, Tokyo, Japan) and elemental distribution maps  were collected in the SEM using energy-dispersive spectroscopy  (Oxford Oxford EDS, FEI). Crystalline phases present  in the  oxide  scales were  characterized using XRD analysis  (X-Pert,  PANalytic, Westborough, MA).  III.  Results and Discussion  A previous study demonstrated that ZrB2 containing 4 mol% WC had improved oxidation resistance compared with nominally pure ZrB2.29 For nominally pure ZrB2, exposure to air at elevated temperatures produced a ZrO2 porous scale that was not protective.13,15 The addition of WC to ZrB2 led to the formation of WO3 along with ZrO2 during oxidation, which reacted to form a W-rich liquid phase at elevated temperatures.  The improved oxidation resistance of ZrB2 containing WC was attributed to densification of the outer ZrO2 scale by liquid phase sintering.29 Cross sections of the oxide scale formed on  ZrB2 containing WC revealed that the ZrO2-based scale appeared to be dense with a fine grain size.29 However,  the  W concentration in the outer scale was lower than the unoxi dized ZrB2, which was presumably caused by volatilization of WO3. As part of the present study, the WC content of the ZrB2 ceramics was varied to control the amount of W-rich phase in  the oxide scale and potentially improve the oxidation resistance  of the ZrB2-based ceramics.  (1)  Solid Solution Formation  The effect of  the concentration of WC on the oxidation resis tance of ZrB2 was studied for WC additions ranging from 4 to 16 mol%. Below its solubility limit, WC is expected to dissolve into the ZrB2 matrix, as was reported in the previous study29 of the effect of WC additions on ZrB2 oxidation, which used an addition of 4 mol% WC. When WC is added in excess of its  solubility limit, WC will dissolve into the matrix until saturation  Table II. Mass Percent Equivalents for Nominal WC Additions used for the Oxidation Study  Fig. 4.  Optical  images of the oxide scales on nominally pure zirconium  diboride (ZrB2) and ZrB2 with a tungsten carbide (WC) addition of 6 mol% oxidized at 16001C for 3 h. The ZrB2 containing WC exhibited a two-layer oxide structure.  is reached and then the excess would be expected to either be  present as dispersed particles  in the ZrB2 matrix or as a new compound. Fig. 1 compares XRD patterns of ZrB2 with WC addition of 4, 8, and 12 mol% that were sintered at 19751C for  2 h. The major peaks in the three XRD patterns can be indexed  to hexagonal ZrB2 (powder diffraction file (PDF) card number 34-0423). As was observed in the previous study, the peaks were  shifted to higher values of two theta compared with pure ZrB2.  WC (mol%)  4  5  6  WC, tungsten carbide.  WO3 (wt%)  6.7  8.4  9.9  Fig. 5.  Measured scale thicknesses for nominally pure zirconium dibo ride (ZrB2) (circles) and ZrB2 ceramics with tungsten carbide (WC) additions of 4 (squares), 5 (diamonds), and 6 (triangles) mol% oxidized at 16001C for various hold times.  \\x0c', 'April 2011  Oxidation of Zirconium Diboride with Tungsten Carbide Additions  1201  Because W (1.4 A˚ ) has a smaller covalent radius than Zr (1.6 A˚ ),32 substitution of W into the ZrB2 lattice is expected to decrease the average unit cell size, shifting the diffraction peaks  to higher angles. The observations indicate that WC formed a  solid solution with ZrB2. In addition, a low-intensity peak appeared around 401 2y in the pattern of ZrB2 with the 8 mol% WC addition. Additional peaks were observed in ZrB2 with 12 mol% WC, which allowed the unknown phase to be indexed  to WB. The WB may be a product of excess WC (i.e., WC that  did not go into solid solution with ZrB2) reacting with B4C and/ or ZrB2, which would result in the formation of nonstoichiometric compounds such as B4\\x00xC and/or ZrB2\\x00x. Based on XRD analysis, the WC solubility limit in ZrB2 appears to be slightly o8 mol%. Therefore, the WC additive concentration must remain below 8 mol% to retain a single phase ceramic.  Another factor in optimizing the amount of the WC additions  is the volume change during oxidation. When W is oxidized to the volume increases by B230%, which means form WO3, oxidation of a small amount of W produces a large volume of  the  WO3. Materials that form protective oxide scales, such as Al, often have substantial volume increases upon oxidation. For Al, the volume increases by B25% when it oxidizes to form a-Al2O3. Tailoring the overall volume increase when W-containing ZrB2 oxidized may also be beneficial for the formation of a dense ZrO2based protective scale. However, if the volume increase is greater  than some limit, and/or the local stress due to volume change is  greater than the strength of the oxide layer, the material may crack  or swell during oxidation, which would compromise the integrity  of the oxide scale. Further, if WC is added in excess of its solubility  limit in ZrB2, the oxidation of discrete WC particles to WO3 could generate local stresses that may lead to fracture of the oxide scale.  In addition,  formation of discrete WO3 particles may lead to longer times for liquid phase formation because of the lower con Fig. 6.  Comparison of nominally pure zirconium diboride (ZrB2) and ZrB2 with tungsten carbide (WC) additions of 4, 5, and 6 mol% after oxidation at 16001C for 6 h.  tact area with ZrO2 compared with the intimate mixing of WO3 and ZrO2 that would be expected to occur due to oxidation of a WC-ZrB2 solid solution. Based on these considerations, WC additions were selected to stay below the WC solubility limit  in  ZrB2, which is about 8 mol%.  (2)  WC Additions  Isothermal oxidation studies were initially conducted at 16001C  for ZrB2 with 0, 4, 8, and 16 mol% WC additions. As seen from the optical microscopy images in Fig. 2, ZrB2 without WC  Fig. 7. Scanning electron microscopic (SEM) image (a) of zirconium diboride (ZrB2) containing 6 mol% tungsten carbide (WC) addition oxidized at 16001C for 3 h along with EDS element maps for O (b), Zr (c), and W (d). The SEM image shows a two-layer oxide scale. The EDS maps show that the  outer scale (farther right in the images) has lower W (d) and O (b) contents than the inner layer (center of images) indicating depletion of WO3 and probably B2O3 from the outer layer.  \\x0c', '1202  Journal of the American Ceramic Society—Zhang et al.  Vol. 94, No. 4  oxidation at 16001C for 1 h, but ZrB2 with 6 mol% WC addition gained only 6 mg/cm2. Presumably, this was due to an increased  evaporation of WO3 from the ceramic with the higher starting WC addition. The difference in weight gain among the three  materials decreased with oxidation time. This may indicate that  the outer ZrO2 times for lower WC additions and that a dense ZrO2 outer layer was not completely formed under these conditions. However,  shorter oxidation  less protective at  layer was  the protective ability of all of  the ZrO2 layers for all different WC additions was similar after 3 h.  (B)  Scale Thickness:  The thickness of the outer scale was  measured to complement weight gain measurements. The crosssections of the scales produced by oxidation at 16001C for 3 h on  nominally pure ZrB2 and on the ZrB2 containing 6 mol% WC additions are compared in Fig. 4. The scale on ZrB2 containing WC had an obvious two-layer structure, with a lighter outer  layer on top of a darker inner layer. For all of oxidized spec imens, thicknesses of the oxidized scales were measured from the  outer surface to the underlying unoxidized ZrB2, thicknesses on all ceramics with WC additions accounted for  i.e. measured  both layers. The scale thicknesses for nominally pure ZrB2 and ZrB2 with 4, 5, and 6 mol% WC additions oxidized at 16001C for 1-6 h are compared in Fig. 5. The scale on nominally pure ZrB2 for up to 6 h was much thicker (B3.2 mm) than the scale on ZrB2 containing 6 mol% WC, which was B0.75 mm after 6 h. Furthermore, for nominally pure ZrB2, the oxide scale thickness increased linearly with time for oxidation at 16001C,  just as weight gain had increased linearly with time, which in dicated that the scale was not protective. In addition, severe edge  cracking was observed for  scale on nominally pure ZrB2 (Fig. 6). In contrast, ZrB2 ceramics containing 4-6 mol% WC  the  addition  exhibited  severe  edge  cracking  during  oxidation,  possibly due to phase transformations  or  the difference  in coefficients of  in the ZrO2 thermal expansion (CTE)  scale and/  between the  and ZrB2. In addition, ZrB2 with 8 16 mol% WC additions had surface cracking, which was pre scale  and  sumably due to the volume change during oxidation or residual  thermal stresses due to a mismatch in CTE between the oxide  layer and the underlying ceramic during cooling. Based on the  proposed WC solubility limit in ZrB2, discussed earlier, and the cracking observed for materials with higher WC additions, WC  additions of 4, 5, and 6 mol% were selected for further study.  (3)  Oxidation Behavior  (A) Weight Gain:  Figure 3 compares the weight gains for  nominally pure ZrB2 with those for ZrB2 with 4, 5, and 6 mol% WC additions oxidized at 16001C for 1, 2, and 3 h. The weight  gain for nominally pure ZrB2 increased linearly with time at 16001C from B20 mg/cm2 after 1 h to B27 mg/cm2 after 3 h. The behavior indicated, again, that the ZrO2 scale formed during oxidation of ZrB2 was not protective. In contrast, ZrB2 containing 4-6 mol% WC had lower weight gains after 1, 2, and 3 h,  compared with nominally pure ZrB2. For example, the weight gain for ZrB2 containing 6 mol% WC was B7 mg/cm2 after 1 h 16001C compared with B20 mg/cm2 at for nominally pure ZrB2. Hence, the addition of WC led to lower weight gains during oxidation at 16001C as compared with nominally pure ZrB2. The weight gain decreased with the increasing oxidation time  for the three ZrB2 ceramics with WC additions. The phenomenon was attributed to the partial volatilization of WO3. The vaporization of B2O3(l) has been studied previously12 and can occur by direct volatilization as described by Reaction 2 or re actions that produce volatile species such as BO2(g), B2O2(g), or B2O(g).  B2O3 ðlÞ ¼ B2O3 ðgÞ  (2)  (l) becomes severe above 12001C  The vaporization of B2O3 and is often considered to be complete after oxidation at temperatures of 15001C or above. Oxidation of ZrB2 to ZrO2 (assuming B2O3 is completely vaporized) produces a net weight gain of 9.2%. When WC is added, a second possible weight loss  process is WO3 vaporization by Reaction 3 or similar reactions that produce dimeric, trimeric, or higher order tungsten oxides.  WO3 ðsÞ ¼ WO3 ðgÞ  (3)  The oxidation of WC to WO3 or another vapor species) produces a weight gain of 18.4%.  (assuming C leaves as CO  Volatilization of WO3(s) could result (i.e., 100% compared with the original mass of WC). For ZrB2 with WC additions, the total mass change will, therefore, be  in complete weight  loss  a  combination  of  the  9.2% weight  gain  due  to  oxidation  of ZrB2 to ZrO2 (assuming complete loss of B2O3 liquid) combined with the mass loss due to the volatilization of part or all of  the WO3(s). An upper limit of mass loss due to WO3 vaporization would be the amount of WC added in terms of mass per cent. The mass percent equivalents of the WC additions are give  in Table II. If all of the WO3(s) were volatilized from ZrB2 containing 6 mol% WC, the mass loss would be approximately  equivalent  to the mass gain due to ZrO2 mass change of approximately zero. Because the overall mass  formation for a net  change is a combination of mass gain due to ZrO2, B2O3, and WO3 formation and mass loss due to volatilization of unknown amounts of B2O3 and WO3, weight gain alone (shown in Fig. 3) does not provide a representative measure of the oxidation of  these materials. The decrease  in weight gain at  longer  times,  shown in Fig. 3,  is likely due to an increasing volatilization of  unknown amounts of WO3 with time. Also shown in Fig. 3, the weight gain of ZrB2 containing WC decreased with increasing WC additions. For example, ZrB2 with 4 mol% WC addition had a 12 mg/cm2 weight gain after  Fig. 8.  Higher magnification scanning electron microscopy (SEM) im age (a) and EDS map (b)  for W near the boundary between the outer  and inner layers in the oxide scale on ZrB2 containing 6 mol% WC.  \\x0c', 'April 2011  Oxidation of Zirconium Diboride with Tungsten Carbide Additions  1203  had much lower scale thicknesses and the growth rates appeared to decrease with isothermal hold time at 16001C. No cracks were  visible in the scales on ZrB2 ceramics containing WC additions of 4, 5, or 6 mol%. Not only did the addition of WC reduce  the outer layer. This may indicate that the inner layer was denser  and had more WO3-rich liquid phase compared with the outer layer. A higher magnification SEM image and EDS map of W  near the boundary between the outer and inner layers (Fig. 8)  mass gain, but it also decreased the scale thickness. Hence, the  indicated that  the outer  layer was  less dense and appeared to  oxide  scales on ZrB2 hanced oxidation protection compared with the scale on nom containing WC additions provided en inally pure ZrO2.  (C)  Scale Morphology:  As seen from Fig. 4, the scale on  nominally pure ZrB2 consisted of a single layer of ZrO2. In contrast, the scales that formed on ZrB2 containing 4, 5, or 6 mol% WC were made up of two layers, a lighter colored outer layer  that appeared similar to the ZrO2 scale on nominally pure ZrB2 and a darker inner layer. X-Ray diffraction analysis of the outer  layer  (not shown)  indicated that all of  the peaks could be in dexed to monoclinic ZrO2 crystalline phases containing W or B were found in the outer  (PDF card number 37-1484). No  layer. Further, no amorphous hump was observed for the outer  layer,  indicating that any amorphous phase was present only in  limited quantities. EDS mapping of ZrB2 containing 6 mol% WC revealed that both layers contained Zr, O, and W, but the  relative concentrations were different (Fig. 7). As seen from the  have a larger grain size with a lower W concentration.  The loss of WO3 from the outer layer of the scale appears to be accompanied by ZrO2 grain growth. Presumably, the initial stage of oxidation produces a dense outer ZrO2 scale that contains a W-rich phase. Owing to the high vapor pressure of tung sten  oxides  at  elevated  temperatures,  the  tungsten  oxides  evaporate from the surface,  leaving behind a less dense outer  oxide scale with larger grain size, compared with the initial stage.  Even with recession of WO3 and microstructure degeneration in the outer ZrO2 layer, the overall oxidation resistance remains better than the porous ZrO2 pure ZrB2. In addition, the inner scale retains WO3 and remains dense, which provides improved oxidation protection compared  scale that develops on nominally  with the ZrO2 scale that develops on nominally pure ZrB2. Optical microscopy (Fig. 9) was used to image microstructures  of cross sections of the scales on nominally pure ZrB2 and ZrB2 with WC additions of 4, 5, and 6 mol% after oxidation at 16001C  EDS mapping,  the outer layer had lower concentrations of W  for 6 h. For the ceramics containing WC additions, the total scale  and O compared with the inner layer. Combining the results of  thicknesses were nearly identical, as indicated in Fig. 5. However,  both the X-ray diffraction and the EDS mapping, the outer layer  the thickness of the inner layer increased as the WC content of  consisted  of mainly monoclinic ZrO2 amount of a second phase that contained W and O that was below the detection limit of XRD (i.e., oB5 wt%). The reduced W content of the outer layer is likely due to partial va (white) with  small  a  porization of tungsten oxide during oxidation at the elevated oxidation temperature (16001C). Because the crystalline phase  was not  identified,  the valence status of W could not be spec ified, hence the W was assumed to be present as WO3 because this is the stable compound indicated on the phase diagram.29  The inner layer had a higher W content compared with the  outer layer. In addition, the inner layer had more Zr and O than  the ceramic increased. Because the overall scale thickness was the  same, the thickness of the outer layer decreased as the WC con tent of the ceramic increased. Based on the thickness of the inner  oxide layer, 6 mol% WC may be the optimum addition in terms  of providing the best oxidation resistance yet still maintaining a  single phase ceramic and a crack-free oxide scale.  Figure 10 compares microstructures of  the cross  sections of  scales on nominally pure ZrB2 and the ZrB2 containing WC additions oxidized at 16001C. Several differences were revealed.  First,  the oxide  scale on ZrB2 denser ZrO2 outer scales that consisted of equiaxed ZrO2 grains  containing WC had  ceramics  Fig. 9. Comparison of cross-section microstructures of nominally pure zirconium diboride (ZrB2) (a) and ZrB2 with tungsten carbide (WC) additions of 4 (b), 5 (c), and 6 mol% (d) after oxidation at 16001C for 6 h.  \\x0c', '1204  Journal of the American Ceramic Society—Zhang et al.  Vol. 94, No. 4  Fig. 10. Comparison of microstructures of zirconium diboride (ZrB2) oxidization outer scales on nominally pure ZrB2 (a) and ZrB2 with 4 (b), 5 (c), and 6 mol% (d) tungsten carbide (WC) additions oxidized at 16001C for 3 h.  (Fig. 10(b-d)). Second, the average grain size of the ZrO2 in the scales increased with the increasing WC content. The increase in  grain size may be due to the formation of a greater amount of  liquid phase in the oxide scale on the ZrB2 ceramics with higher WC additions during oxidation. The larger liquid phase content  could allow for more rapid coarsening of  the ZrO2 grains. contrast to the equiaxed microstructure of the scales on ceramics  In  containing WC,  the nominally pure ZrB2 had a columnar, porous ZrO2 structure (Fig. 10(a)) that did not provide significant oxidation protection.  IV.  Conclusion  The addition of a W compound, WC in this case,  improved the  oxidation resistance of ZrB2. During oxidation, the presence of WO3 in the oxide scale resulted in liquid phase sintering of ZrO2, which modified the microstructure of the scale and was thought  to increase its relative density. As a result, WC-containing ZrB2 had improved oxidation resistance compared with nominally  pure ZrB2 as measured by both mass gain and oxide scale thickness. Based on the theoretical considerations and analysis of  experimental  results,  the  following conclusion can be drawn  from this study: solubility of WC in ZrB2 wasoB8 mol%. For additions below this limit, WC formed a solid solution with  The  (1)  ZrB2 during densification. The ZrO2 oxidation scale for ZrB2 with \\x158 mol% WC (2) exhibited cracking, which may be due to the volume increase (B230%) associated with the oxidation of W to WO3. Owing to the cracking, the amount of WC added to ZrB2 was limited to o8 mol% for the detailed oxidation studies. (3) During oxidation, WC-containing ZrB2 oxidized to form B2O3, WO3, and ZrO2. A liquid phase formed due to the presence of WO3 and ZrO2, which resulted in ZrO2 grains with an equiaxed microstructure. In contrast, the ZrO2 scale that formed on nominally pure ZrB2 was porous and consisted of columnar grains.  (4) Weight gain alone was not a reliable measure of the ox idation resistance due to the partial vaporization of B2O3 and WO3. Therefore, the thickness of oxide scale was measured to evaluate the oxidation resistance of ZrB2 without WC additions. The scale on nominally pure ZrB2 ceramics was much thicker (43 mm after 6 h at 16001C) than the scale that formed on ZrB2 ceramics containing WC (B0.75 mm after 6 h at 16001C).  ceramics with and  (5)  Owing to evaporation of WO3 at the oxidation temperature, WC-containing ZrB2 ceramics developed a two-layer scale structure, which consisted of a porous ZrO2 outer layer and a denser inner layer containing ZrO2 with more WO3. In contrast, nominally pure ZrB2 had a single layer of highly porous ZrO2 with a columnar microstructure, which did not protect the un derlying ceramic from oxidation. For WC-containing ceramics,  both the scale thicknesses and mass gains were lower compared  with nominally pure ZrB2. (6) Considering the volume change associated with reaction  and the solubility limit for W in ZrB2 as well as the evaporation of WO3 from the outer scale layer, the addition of 6 mol% WC appears to be the optimum concentration. It may be possible to  further improve the oxidation resistance of ZrB2 by identifying other additives that promote liquid phase sintering of the oxide  scale, but that have lower vapor pressures than WO3 at the oxidation temperature, which could reduce the overall scale thick ness by stabilizing the  inner  layer  that  is  thought  to provide  oxidation protection.  References  1M. J. Gasch, D. T. Ellerby, and S. M. Johnson,  ‘‘Ultra High Temperature  Ceramic Composites’’; pp. 197-224 in Handbook of Ceramic Composites, Edited  by N. P. Bansal. Kluwer Academic Publishers, Boston, 2005. 2W. G. Fahrenholtz, G. E. Hilmas,  I. G. Talmy,  and  J. A. Zaykoski,  ‘‘Refractory Diborides of Zirconium and Hafnium,’’ J. Am. Ceram. 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The American Ceramic Society,  Westerville, OH, 1998. 31S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Pressureless Densificat ion of Zirconium Diboride with Boron Carbide Additions,’’ J. Am. Ceram. Soc.,  89 [5] 1544-50 (2006). 32F-175,  ‘‘Crystal and Ionic Radii of  the Elements’’;  in CRC Handbook of  Chemistry and Physics, 62nd ed, CRC Press, Boca Raton, FL, 1981.  &  \\x0c']"
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  "_id": 175,
  "PDF": "Oxidation of Zirconium Diboride–Silicon Carbide at 15001C at a Low Partial Pressure of Oxygen.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 89 [10] 3240 - 3245 (2006)  DOI: 10.1111/j.1551-2916.2006.01229.x  r 2006 The American Ceramic Society  Oxidation of Zirconium Diboride-Silicon Carbide at 15001C at a Low  Partial Pressure of Oxygen  Alireza Rezaie,* William G. Fahrenholtz,*,w  and Gregory E. Hilmas*  Department of Materials Science and Engineering, University of Missouri Rolla, Rolla, Missouri 65409  The  oxidation  behavior  of  zirconium diboride  containing  30  vol% silicon carbide particulates was investigated under reducing conditions. A gas mixture of CO and B350 ppm CO2 was used to produce an oxygen partial pressure of B10\\x0010 Pa at 15001C. The kinetics of the growth of the reaction layer were examined for reaction times of up to 8 h. Microstructures and  chemistries of reaction layers were characterized using scanning  electron microscopy and X-ray diffraction analysis. The kinetic  measurements,  the microstructure analysis, and a thermody namic model  indicate that oxidation in CO-CO2 produced a non-protective oxide surface scale.  I.  Introduction  ZrB2 exhibits rapid linear oxidation kinetics. At atures, the rate of vaporization of B2O3 is high compared with its rate of formation. The removal of the B2O3 layer leaves behind a porous ZrO2 scale, which does not provide passive oxidation protection to the underlying ZrB2.18 The oxidation behavior of ZrB2 has also been studied at reduced oxygen pressures. At a total pressure of 33, 429 Pa of pure  these temper oxygen, the kinetics of the oxidation were influenced by vapor ization of B2O3 similar to what was described above for oxidation in air.16 In another study, the oxidation of ZrB2 at 10561C at various partial pressures of oxygen (pO2 5 13, 169-99, 247 Pa) showed that the oxidation rate increased with increasing oxygen partial pressure.15 The authors concluded that the reaction rate was limited by a diffusion mechanism.15 In a similar study, no  in materials  RECENT interest (TPS) for hypersonic aerospace vehicles and reusable atmospheric re-entry vehicles has resulted in significant research activity focused on ultrahigh-temperature ceramics (UHTCs).1-12  for thermal protection systems  UHTCs are a group of materials,  including TaC, ZrB2, ZrC, HfB2, HfC, and HfN, that can potentially be used at temperatures above 15001C. The characteristics of UHTCs, such as a high melting temperature (430001C), strength at high temper atures, and oxidation resistance, allow them to be used in harsh  environments associated with hypersonic flight, atmospheric re-entry, and rocket propulsion.1,4 Among the UHTCs, ZrB2 has the lowest theoretical density (6.09 g/cm3), which can be an for aerospace applications.13  advantage over other candidates  Additionally, ZrB2 has excellent resistance to thermal shock due to its high thermal conductivity (65-135 W \\x01 (m \\x01 K)\\x001).13 The oxidation of ZrB2 has been studied by a number of investigators.14-17 When ZrB2 is exposed to air at elevated temperatures, ZrO2 and B2O3 (l) are formed according to the following reaction:  ZrB2þ 5 2  O2 ðgÞ ! ZrO2þB2O3 ðl Þ  (1)  The B2O3 is present as a liquid above its melting temperature (B4501C), while the ZrO2 tends to form a porous crystalline solid. Below 11001C, passive oxidation behavior with parabolic kinetics has been observed.15 Oxidation protection has been at tributed to the presence of a continuous layer of B2O3, which acts as an oxygen diffusion barrier. Above 11001C, evaporation  of B2O3 reduces the effectiveness of ZrO2 does not offer effective protection to the underlying ZrB2.16 Between 11001 and 14001C, para-linear kinetics have been observed.16 In para-linear kinetics, the overall rate of mass  the diffusion barrier. The  change is a combination of weight gain due to oxide formation loss due to B2O3 volatilization.8,16 Above 14001C,  and weight  N. Jacobson—contributing editor  Manuscript No. 21636. Recevied March 27, 2006; approved June 7, 2006.  This material is based upon work supported by the National Science Foundation under  grant number DMR-0346800.  *Member, American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: billf@umr.edu  relationship between the oxidation rate and oxygen partial pressure was determined at 10001C,17 which was Kuriakose and Margrave’s15 conclusion that  the rate was di in contrast with  rectly proportional to the partial pressure of oxygen when the total pressure was 1.013 \\x02 105 Pa (1 atm). In this case, suggested that the diffusion of molecular oxygen was the ratecontrolling step.17  it was  Attempts have been made to improve the oxidation resistance of ZrB2-based materials through appropriate additives.19-22 For temperatures above 12001C, the addition of SiC provides im proved oxidation resistance by encouraging the formation of a  borosilicate glass layer on exposed surfaces, which results in reduced oxygen transport.1,21,23,24 The SiC phase begins oxidizing appreciably between 11001 and 13001C, resulting in the forma tion of a silica-rich glassy layer (Reaction (2)) that imparts sig nificant improvement to the oxidation resistance of the diboride  at higher temperatures:  SiCþ 3 2  O2 ðgÞ ! SiO2 ðl Þ þ COðgÞ  (2)  The SiO2-containing scale layer remains protective up to at least 15001C as SiO2 is significantly less volatile than B2O3 (by a factor of B105) at these temperatures.23 Thus, ZrB2-SiC exhibits passive oxidation behavior with parabolic kinetics over a  much greater temperature range than has been reported for pure  ZrB2. For example, the addition of SiC has been reported to decrease the normalized weight gain upon heating to 15001C in air without an isothermal hold from 13.0 mg/cm2 for pure ZrB2 to 1.8 mg/cm2 for ZrB2 containing 30 vol% SiC.25 The oxidation behavior of ZrB2-SiC at reduced oxygen partial pressures has not been reported. Exposure of ZrB2-SiC to air at 15001C results in the formation of a layered surface structure. As discussed above, the outer  surface of the oxidized specimen is a silica-rich glassy layer that  provides the passive oxidation protection with parabolic kinet ics. Several investigators have described the formation of a ‘‘SiC-depleted layer’’ beneath the outer oxide layer.6,26-30 The  SiC-depleted layer has been reported to consist of ZrB2 or ZrO2, depending on the conditions of the experiment. The structure of  this layer is similar to the original ZrB2-SiC, but with the SiC partially or fully removed. Below the SiC-depleted layer is the  unaffected ZrB2-SiC base material. The development of the layered structure during ZrB2-SiC oxidation in air at 15001C was  3240  \\x0c', 'analyzed with the aid of a thermodynamic model that involved volatility diagrams for ZrB2 and SiC.18,28 The model suggested that the high vapor pressure of SiO (g) under reducing condi tions led to the active oxidation of SiC (Reaction (3)) beneath  the SiO2 layer, region26,28:  resulting in the formation of  the SiC-depleted  SiC þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  (3)  The purpose of  this article is  to describe the oxidation be havior of ZrB2-SiC under low pO2). The composition of the reaction layer formed on the surface was investigated along with the kinetics of oxidation at 15001C in a  reducing conditions  (i.e.  CO-CO2 mixture. The results will be analyzed and compared with the conditions thought to be responsible for the formation  of the SiC-depleted layer during oxidation of ZrB2-SiC in air.  II.  Experimental Procedure  (1)  Processing  Commercially available ZrB2 MA) was used to prepare the materials for this study. This powder had a reported purity of 499% (metals basis) and a reported average particle size of 2 mm. The SiC powder (Grade UF-10, H.C. Starck) was predominantly a-SiC. It had a reported purity of 98.5% and a reported average particle size of 0.7 mm. Batches  (Grade B, H.C. Starck, Newton,  containing 70 vol% ZrB2 and 30 vol% SiC were prepared. The powders were attrition milled (Model 01-HD, Union Process,  Akron, OH) to reduce particle size and promote intimate mixing  of the ZrB2 and SiC. A 750 mL fluoropolymer-coated bucket was charged with B250 mL hexane, B150 g of powder, and B3000 g of ZrO2 milling media (B3.5 mm diameter milling was performed at 600 rpm for 2 h. The  spheres). Attrition  solvent was  removed using rotary evaporation (Model Rotavapor R-124, Buchi, Flawil, Germany) at a temperature of 701C, a vacuum of 200 mmHg (B27 kPa), and a rotation speed of 150 rpm.  Powders were densified by hot pressing (Model HP-3060, Thermal Technology, Santa Rosa, CA) at 19501C for 45 min at  a pressure of 32 MPa,  in a graphite die lined with graphite foil  that was coated with BN. The furnace was heated using a heating schedule that has been reported previously.31 The furnace was heated at a rate of 101C/min under vacuum (B20 Pa) up 16501C,  to  at which time the atmosphere argon. Above B8001C,  was  switched  to flowing  an infrared thermometer  (Model OS 3708, Omega Engineering, Stanford, CT) was used  to monitor the temperature of the graphite die. When the die temperature reached 19501C, a uniaxial applied. After 45 min at 19501C, the furnace was cooled at B201C/min to room temperature. The load was removed when the die temperature was below 17501C. Billets with a diameter of B40 mm and a thickness of B5 mm were produced. Specimens with dimensions of 6 mm \\x02 4 mm \\x02 4 mm were diced from the billets for oxidation.  load of 32 MPa was  (2)  Oxidation  The experimental portion of this study focused on exposing the ZrB2-SiC specimens to CO containing B350 ppm CO2, which produced an oxygen partial pressure of B10\\x0010 Pa at 15001C. This oxygen partial pressure was selected based on the infor mation obtained from a thermodynamic model  that predicted  an oxygen partial pressure in this range in the SiC-depleted zone of ZrB2-SiC.28 A MoSi2 furnace (Model 0000543 Rapid Temperature  during  oxidation  resistance-heated  horizontal  tube  Furnace, CM Inc., Bloomfield, NJ) equipped with a high-purity  alumina tube with a diameter of 6.35 cm was used for  the  oxidation studies. Before oxidation, specimens were cleaned in  an ultrasonic bath with acetone. Specimens were placed on an  alumina plate,  that was on a D-tube,  inserted into the center  of  the furnace and leveled. The ends of  the tube were sealed  with gas-tight end caps. A flowing atmosphere was maintained.  Oxidation studies were conducted in a mixture of CO containing 36177 ppm CO2. A flow rate of B100 cm3/min was used. Specimens were heated at B51C/min to 15001C and held for  0, 0.5, 1, 2, 4, and 8 h. The weight of each specimen was meas ured before and after oxidation to determine the weight change as a function of exposure time at 15001C. The thicknesses of the  resulting reaction layers were measured from polished cross-sec tions. An image analysis program (Image J, U.S. National In stitutes of Health, Bethesda, MD) was used to determine layer  thickness and standard deviation from a minimum of 20 meas urements per specimen. For comparison, an additional specimen was oxidized in flowing air at a rate of B1.8 L/min at 15001C  with a hold time of 30 min.  (3)  Characterization  The microstructure of oxidized specimens was characterized us ing scanning electron microscopy (SEM; S-570, Hitachi, Tokyo,  Japan) along with energy-dispersive spectroscopy (EDS; AAT,  X-ray Optics, Gainesville, FL)  for chemical analysis. Charac teristic X-ray mapping at an accelerating voltage of 5 kV was  used to analyze  the distribution of different  elements  in the  specimens. Samples were prepared for microscopy by cutting cross-sections and then polishing to a 0.25 mm finish using dia mond abrasives. Specimens were coated with carbon after pol ishing and before imaging. In addition to SEM/EDS, specimens  were  examined  using  grazing  incidence  X-ray  diffraction  (GXRD; X’Pert MRD, Panalytical, Almelo,  the Netherlands)  to determine the reaction products. Secondary ion mass spec troscopy (SIMS; PHI TRIFT III, Physical  Instrument, Eden  Prairie, MN) was used to determine the elemental composition  of the protective external scale in the specimen exposed to air.  III.  Results and Discussion  The measured bulk densities of the hot-pressed billets ranged from 5.13 to 5.18 g/cm3. Using a simple rule of mixtures calculation that assumed true densities of 6.09 g/cm3 for ZrB2 and 3.21 g/cm3 for SiC, the theoretical density of ZrB2 containing 30 vol% SiC should be 5.23 g/cm3. All of the specimens were hot pressed to near theoretical density (498%), with no indication  of open porosity. Based on the high relative density and the lack  of any open porosity, porosity should not have a significant  effect on the oxidation behavior.  An example of the microstructure of the ZrB2-SiC specimens used in the oxidation experiments in this study is presented in  Fig. 1. The dark phase is SiC, which is distributed in the lighter colored ZrB2 phase. The SiC particulates appear to be uniformly distributed in the ZrB2 matrix. No indication of porosity was observed by SEM, which was consistent with the density meas urements. Based on analysis of the microstructures of ZrB2-SiC  20 µµm  SiC  Fig. 1. Scanning electron micrograph of a polished, thermally etched ZrB2-SiC specimen prepared by hot pressing at 19501C for 45 min.  October 2006  Oxidation of Zirconium Diboride-Silicon Carbide  3241  \\x0c', 'prepared using an identical procedure, the ZrB2 phase had an approximate grain size of 2.5 mm and SiC had an approximate grain size of 1.7 mm.32  (1)  High pO2 Regime (Air)  Oxidation of ZrB2-SiC at 15001C for 30 min in air (pO2 5 20 260 Pa) produced a structure that consisted of four layers: (o10 mm), (2) a Zr-rich oxidized layer embedded in amorphous SiO2 (o3 mm), SiC-depleted layer (note that SiC is only partially depleted and that may be ZrO2 or ZrB2 (B10 mm), and (4) unaffected ZrB2-SiC. These layers are shown in Fig. 2. The formation of this four-layer structure is consistent with obser (1) a continuous  surface oxide layer  (3) a  not  fully removed)  vations of other investigators who oxidation.6,23,27,29 The formation of  have  studied ZrB2-SiC the layered structure was  also described by a thermodynamic model that used volatility diagrams for ZrB2 and SiC.28 A similar volatility diagram has been produced for the present study and is shown in Fig. 3.  Based on the thermodynamic model, exposure of ZrB2-SiC to air (log pO2 5 4.3) at 15001C should produce SiO2, ZrO2, and B2O3. Typically, a silica-rich glassy layer is observed on the outer surface. The B2O3 is depleted from the borosilicate glass due to preferential evaporation because of its high vapor pressure  relative to SiO2. Analysis of the composition across the interface between the outer oxide layer and the SiC-depleted layer by  SIMS (Fig. 4) confirmed that the surface layer was boron defi cient compared with the underlying layers after oxidation at 15001C. The boron concentration showed a sharp increase at the  interface. The outer oxide  layer  contained almost no boron,  while the SiC-depleted layer contained almost no oxygen. The  silica-rich surface layer is thought to act as a diffusion barrier for  oxygen, resulting in passive oxidation behavior with parabolic weight gain kinetics.29 A weight gain of 24% is expected based  on the stoichiometry of Reactions (1) and (2), which show that  the weight of SiO2 and ZrO2 formed is more than the sum of weights of ZrB2 and SiC consumed, plus B2O3 and CO lost as gaseous species (Table I). Experimentally, weight gain was obthe specimen exposed to air at 15001C for 30 min served for (B1.3 mg/cm2).  Beneath the silica-rich surface layer,  the partial pressure of  oxygen has been predicted to be significantly lower than the ex ternal atmosphere due to the oxygen chemical potential gradient that develops across the diffusion barrier.28 In fact, the oxygen  partial pressure is low enough to promote the active oxidation of  SiC beneath the scale to form SiO (g) and CO (g), presumably  by Reaction (3). The active oxidation of SiC beneath the outer  oxide layer is thought to produce the SiC-depleted layer that has been observed for the oxidation of ZrB2-SiC in air.6,27-29 Even though it is thought that SiC undergoes active oxidation  (presumably linear weight loss kinetics) as part of the oxidation  of ZrB2-SiC, the overall oxidation process in air is expected to show parabolic weight gain kinetics (passive oxidation, protec tive scale) due to the stability of been reported previously.29  the silica surface layer as has  (2)  \\x0015 \\x15 pO2 \\x15 10 \\x005 Pa) Intermediate pO2 Regime (10  From the thermodynamic model constructed to understand ox idation of ZrB2-SiC in air, exposure of ZrB2-SiC to intermediate oxygen partial pressures should result in active oxidation of  SiC. In this pO2 regime, tively high, which does not allow a protective oxide layer  the vapor pressure of SiO (g)  is rela to  form on the surface. Depending on the precise pO2, the SiC will either be removed by active oxidation (below pO2B8.8 \\x02 10\\x0013 will then volatilize (above pO2B8.8 \\x02 10\\x0013 Pa) (Reaction (4)). Pa) (Reaction (3)) or it will oxidize to SiO2 (Reaction (2)), which Likewise, the ZrB2 will either be stable below pO2B1.9 \\x02 10\\x0011 Pa or will oxidize to form ZrO2 and B2O3. If B2O3 does form at this temperature and in this pO2 regime, it should volatilize, leaving ZrO2 as the only stable condensed phase. To date, no experiments have reported oxidation of ZrB2-SiC in this oxygen partial pressure regime.  SiO2 ðl Þ ! SiOðgÞ þ 1 2  O2 ðgÞ  (4)  Specimens of ZrB2-SiC were exposed to a CO-CO2 mixture with pO2 B10\\x0010 Pa at 15001C. Whereas a dense, stable SiO2 surface layer formed when ZrB2-SiC was exposed to air, exposure to the low pO2 (pO2 \\x19 10\\x0010Pa) produced a porous surface layer. No protective layer was observed using SEM. Figure 5  10 µµm  Fig. 2. Scanning electron micrograph showing the layered structure that develops on ZrB2-SiC when it is exposed to air at 15001C for 30 min.  0  5  10  − 20 − 30  − 25  − 20  − 15  − 15  − 10  − 10  − 5  − 5  0  5  10  Log pO2 (Pa)  g o L  P  a  r  t  i  a  l  P  r  s s e  u  r  e  (  P  a  )  ZrB2 (cr) + SiC (cr)  ZrB2 (cr) + SiO2 (I)  SiO (g)  SiO (g)  SiO2 (g)  B(g)  BO (g)  B2O2 (g)  B2O3 (g)  BO2 (g)  Fig. 3. Volatility diagram for diagram for ZrB2 at 15001C.28  SiC superimposed  on  the  volatility  B  (  c  n u o  t  s  /  s  )  Depth (nm)  O  (  c  n u o  t  s  /  s  )  Fig. 4.  Variation in the concentrations of boron and oxygen across the  interface between the outer oxide layer and the SiC-depleted layer.  3242  Journal of the American Ceramic Society—Rezaie et al.  Vol. 89, No. 10            \\x0c', 'shows a SEM micrograph and three compositional maps of a  ZrB2-SiC specimen that was exposed to the CO-CO2 mixture at 15001C for 30 min. Figure 5(a) shows a porous layer on the  specimen surface. No sign of  a  continuous protective  layer,  which could impart diffusion-controlled oxidation protection  to the underlying material, was observed. The white material  observed on the specimen surface (Fig. 5(a))  is contamination  from polishing the porous  specimen. Composition maps pro duced using EDS analysis (Fig. 5) showed that zirconium and  oxygen were present  in the reaction layer, but  that silicon was  absent, suggesting that  the reaction layer was ZrO2. Additionally, XRD analysis revealed that monoclinic ZrO2 was the only crystalline phase in this layer (data not shown, but indexed to  ICDD card number 86-1445).  The variation of the thickness of the reaction layer as a function of holding time at 15001C in an atmosphere of CO-CO2 is shown in Fig. 6. The thickness of the reaction layer increased from 1673 to 11677 mm when the exposure time increased from 30 min to 8 h at 15001C. From this plot, the thickness of  the reaction layer appeared to increase linearly with time. Each  point on the graph is an average of at least 20 measurements in  different regions of a specimen. The reaction layer thickness was  relatively uniform in different areas of the specimen. For exam ple,  in the specimen with an exposure time of 8 h, the measured reaction layer thickness ranged from 105 to 125 mm, with an average of 116 mm and a standard deviation of 7 mm. The uni formity of the reaction layer resulted in relatively small standard  deviations (6%-18%) for all of the measurements.  The variation of the thickness of the reaction layer with time  can be an indicator of the oxidation mechanism. The linear trend (Fig. 6) suggests reaction rate-controlled kinetics.33 In this  case, the porous ZrO2 layer on the surface did not act as a barrier for diffusion of oxygen to the underlying ZrB2-SiC but it allowed direct flow of the external atmosphere to the reaction  interface. In this case, the ZrB2-SiC exhibited passive oxidation behavior (a stable ZrO2 layer was formed), but the scale was non-protective, resulting in linear  reaction kinetics. Based on  this observation,  it appears that  the rate of  increase in the re action layer  thickness was  limited by the reaction rate at  the  surface of the ZrB2-SiC. The growth rate of the reaction layer was calculated to be B13.6 mm/h.  To further  substantiate  the analysis based on the  reaction  layer thickness increase as a function of time, the weight change  was also measured as a function of time for the exposure of ZrB2-SiC to CO-CO2 at 15001C (Fig. 7). The weight change results are consistent with the trend of the reaction layer thick ness. The plot shows that the weight loss increased linearly with  holding time, suggesting reaction-controlled kinetics in which no  effective protection was provided by the porous ZrO2 layer. The rate of mass loss for ZrB2-SiC in CO-CO2 (pO2B10\\x0010 Pa) was calculated from the data in Fig. 7 and was 3.8 mg (cm2 \\x01 h) \\x001. The results of weight loss measurements were also consistent  with the  volatility diagram, which predicted vaporization of  both SiO2 and B2O3 of oxygen. Weight loss was expected based on the stoichiometry  (l)  (l)  at  intermediate partial pressures  of Reaction (5), which showed that the weight of ZrO2 formed was less than the sum of weights of ZrB2 and SiC consumed,  a)  b)   c)   O   Zr   SE   10 µm  10 µm  10 µm  d)    Si                                                      20   ZrB2-SiC  ZrB2-SiC   Si  10 µm  ZrO2  ZrO2  ZrB2-SiC  ZrO2  ZrB2-SiC  ZrO2  Fig. 5.  Cross-section of ZrB2-SiC heated to 15001C for 30 min in CO- scanning electron microscope image and compositional maps  CO2 for (b) Zr, (c) O, and (d) Si.  (a)  Table I.  Comparison of the Weights of the Reactants and  Reaction Products for ZrB2-SiC Oxidation  Component  Molar weight  (g/mol)  Molar weight  (g/mol)  Weight  change (%)  Reactants  ZrB2 SiC  113  153  40  Products (air)  ZrO2 SiO2 (l) B2O3 (g)  123  183  124  60  70  (vaporized)  Products  (intermediate  pO2)  ZrO2 SiO2 (g)  123  123  \\x0034  60  (vaporized)  B2O3 (g)  70  (vaporized)  October 2006  Oxidation of Zirconium Diboride-Silicon Carbide  3243  \\x0c', 'assuming that SiO2 (l) and B2O3 (l) were lost due to volatilization (Table I):  ZrB2þSiC þ 4O2 ðgÞ ! ZrO2 þ B2O3 ðgÞ þ SiO2 ðgÞ  þ COðgÞ  (5)  (3)  Low-pO2 Regime (pO2r10 \\x0015 Pa)  One additional major pO2 regime is apparent from Fig. 3. At the lowest values of oxygen partial pressure (pO2o10 \\x0015 Pa), both ZrB2 and SiC are stable condensed phases. At the upper end of this pO2 regime, ZrB2 and SiC may both undergo active oxidation as BO (g) and SiO (g) have relatively high vapor pressures.  However, as pO2 decreases, the stability of the condensed phases increases. As with the intermediate-pO2 regime, no experiments have been reported for the oxidation of ZrB2-SiC in this oxygen partial pressure regime. In this case, the experiments are com plicated by the need for more complex gas mixtures. Below about pO2 5 10\\x0015 Pa, CO-CO2 mixtures cannot be used as this is below the limit where CO dissociates to form solid carbon and  CO2, the so-called sooting limit. Thus, it is not possible to produce oxygen partial pressures less than B10\\x0015 Pa using CO- CO2 at 15001C. Use of H2/H2O mixtures that could produce oxygen partial pressures a low pO2 regime at 15001C was not  considered appropriate due to acceleration of B2O3 (l) vaporization in the presence of water vapor. Other techniques such as  metal-metal oxide combinations may also provide pO2 less than \\x0015 Pa at 15001C. 10  IV.  Summary  Specimens of ZrB2 containing 30 vol% SiC were oxidized in a CO-CO2 mixture. The gas contained B350 ppm CO2, resulting \\x0010 Pa at 15001C. A reacin an oxygen partial pressure of B10 tion layer composed of porous ZrO2 was formed on the surface during exposure. The results of weight change experiments and  the variation of the reaction layer thickness with time verified the oxidation of ZrB2-SiC at pO2B10 \\x0010 Pa can be described with linear weight gain kinetics. The results indicate that  that  no protective oxide layer  formed on ZrB2-SiC at partial pressure. The experimental results are consistent with the  this oxygen  predictions of a thermodynamic model, developed to describe  the formation of a SiC-depleted region in ZrB2-SiC, which indicated that the SiC should undergo active oxidation at this pO2. The rate of mass loss for a ZrB2 specimen containing 30 vol% SiC in CO-CO2 at 15001C was 3.8 mg(cm2 \\x01 h)\\x001. Under the same conditions, the average growth rate of the reaction layer was 13.6 mm/h.  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Cutler,  ‘‘Engineering Properties of Borides’’; pp. 787-803 in Ceramics  and Glasses, Engineered Materials Handbook, Vol. 4, Edited by S. J. Schneider Jr.  ASM International, Materials Park, OH, 1991. 14J. B. Berkowitz-Mattuck,  ‘‘High-Temperature Oxidation: III. Zirconium and  Hafnium Diborides,’’ J. Electrochem. Soc., 113 [9] 908-14 (1966). 15A. K. Kuriakose and J. L. Margrave,  ‘‘The Oxidation Kinetics of Zirconium  Diboride and Zirconium Carbide at High Temperatures,’’ J. Electrochem. Soc.,  111 [7] 827-31 (1964). 16W. C. Tripp and H. C. Graham, ‘‘Thermogravimetric Study of the Oxidation of ZrB2 in the Temperature Range of 8001-15001C,’’ J. Electrochem. Soc., 118 [7] 1195-9 (1968).  0  20  40  60  80  100  120  140  0  4  8  y = 12.977 + 13.612x R2 = 0.9569  R  a e  c  t  i  n o  a L  y  e  r  (  µ  m  )  Holding Time (hrs)  6  2  Fig. 6. Variation of the reaction layer thickness with time for exposure of ZrB2-SiC to CO-CO2 at 15001C.  − 3  − 2.5  − 2  − 1.5  − 1  − 0.5  0  0  2  4  6  8  y = − 0.005837 − 0.38186x R2 = 0.97933   W  e  i  h g  t  C  e g n a h  (  m  g  /  c  m  2  )  Holding Time (hrs)  Fig. 7. Weight change as a function of time for exposure of ZrB2-SiC to CO-CO2 at 15001C.  3244  Journal of the American Ceramic Society—Rezaie et al.  Vol. 89, No. 10          \\x0c', 'October 2006  Oxidation of Zirconium Diboride-Silicon Carbide  3245  17R. J. Irving and I. G. Worsley,  ‘‘Oxidation of Titanium Diboride and Zirco ature Monolithic and Fibrous Monolithic Ceramics,’’ J. Mater. Sci., 39, 5951-7  nium Diboride at High Temperatures,’’ J. Less-Common Metals, 16 [2] 102-12  (1968). 18W. G. Fahrenholtz,  ‘‘The ZrB2 Volatility Diagram,’’ J. Am. Ceram. Soc., 88  [12] 3509-12 (2005). 19K. Bundschuh, M. Schuze, C. Muller, P. Greil, and W. Heider, ‘‘Selection of for Use at Temperatures above 15001C in Oxidizing Atmospheres,’’  Materials  J. Eur. Ceram. Soc., 18, 2389-91 (1998). 20G. A. Pankov, G. A. Fomina, D. A. Ivanov, and G. E. Val’yano,  ‘‘Strength  and Scaling Resistance of a Composite Based on Zirconium Diboride,’’ Refrac tories, 35 [9-10] 298-300 (1994). 21K. Kobayashi, H. Sano, K. Maeda, and Y. Uchiyama,  ‘‘Oxidation Behavior  of Graphite-B4C/SiC/ZrB2 Composite in Dry and Moist Atmosphere,’’ J. Ceram. Soc. Jpn. Int. Ed., 100, 407-11 (1992). 22X. Zhong and H. Zhao,  ‘‘High Temperature Properties of Refractory Com posites,’’ Am. Ceram. Soc. Bull., 78, 98-101 (1999). 23W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of an SiC Addition on  the Oxidation of ZrB2,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 24M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey,  ‘‘Mechanical, Thermal, and Oxidation Properties of Hafnium and Zirconium  Compounds,’’ J. Eur. Ceram. Soc., 19, 2405-14 (1999). 25W. G. Fahrenholtz, G. E. Hilmas, A. L. Chamberlain, and J. W. Zimmer mann,  ‘‘Processing  and Characterization of ZrB2-Based Ultra High Temper (2004). 26E. Opila, S. Levine, and J. Lorincz, ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39 [19]  5969-77 (2004). 27M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties, and Arc Jet Oxidation of Hafnium Diboride/Silicon Car bide Ultra High Temperature Ceramics,’’ J. Mater. Sci., 39 [19] 5925-37 (2004). 28W. G. Fahrenholtz,  of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., accepted. 29E. J. Opila and M. C. Halbig,  ‘‘Oxidation of ZrB2-SiC,’’ Ceram. Eng. Sci.  ‘‘Thermodynamic Analysis  Proc., 22 [3] 221-8 (2001). 30E. V. Clougherty, R. L. Pober, and L. Kaufman,  ‘‘Synthesis of Oxidation  Resistance Metal Diboride Composites,’’ Trans. Metall. Soc., AIME, 242, 1077-82  (1968). 31A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  ‘‘High  Strength ZrB2-Based Ceramics,’’  J. Am. Ceram. Soc.,  87  [6]  1170-2  (2004). 32A. R. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Effect of Hot Pressing  Time and Temperature on the Microstructure and Mechanical Properties of ZrB2- SiC,’’ J. Mater. Sci., accepted. 33W. E. Lee  ‘‘Melt Corrosion of Oxide  and Oxide-Carbon  and S. Zhang,  Refractories,’’ Int. Mater. Rev., 44 [3] 77-104 (1999).  &  \\x0c']"
},{
  "_id": 176,
  "PDF": "Oxidation of zirconium diboride–silicon carbide ceramics under an oxygen partial pressure of 200Pa- Formation of zircon.pdf",
  "Text": "['Corrosion Science 52 (2010) 3297-3303  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200 Pa: Formation of zircon  Dong Gao a, Yue Zhang a,*,  Jingying Fu a, Chunlai Xu b, Yang Song b, Xiaobin Shi b  a Key Laboratory of Aerospace Materials and Performance (Ministry of Education), School of Materials Science and Engineering, Beihang University, Beijing 100191, PR China b National Key Laboratory of Advanced Functional Composite Materials Technology, Beijing 100076, PR China  a r t i c l e  i n f o  a b s t r a c t  The formation and evolution of zircon during oxidation of ZrB2-20 vol.% SiC ceramics under a low oxygen partial pressure of 200 Pa is studied. The formation mechanism of zircon is proposed according to experimental results and thermodynamic consideration. And the main reason to the formation of zircon can be attributed to the active oxidation of SiC. Two steps can be divided for the formation and evolution of zircon: (1) nucleation from silica glass; and (2) crystal growth into prism like particles. Furthermore, the emergence of zircon significantly improves the oxidation resistance performance. Ó 2010 Elsevier Ltd. All rights reserved.  Article history:  Received 4 February 2010 Accepted 2 June 2010 Available online 15 June 2010  Keywords:  A. Ceramic matrix composites B. EPMA C. Oxidation  1. Introduction  Ceramic of borides, such as zirconium diboride (ZrB2) and hafnium diboride (HfB2) are known as ultra-high temperature ceramics due to their extremely high melting points [1,2]. They are considered as potential candidate material of thermal protective system (TPS) for re-entry aircrafts and supersonic vehicles because of their unique combination of high melting points, good thermal shock resistance, and excellent ablation/oxidation resistance [3-6]. Of those UHTCs, ZrB2 has the lowest theoretical density (6.09 g/cm3), which makes it as the most promising material used for the TPS. In addition, the thermal shock resistance property of the ZrB2 ceramics is superior to other ultra-high temperature ceramics due to its higher thermal conductivity (65-135 W/(m \\x02 K)\\x001) [7,8]. Oxidation resistance performance is the major challenge in the development of ZrB2 ceramics which severely undermines their ability to survive oxidizing condition at high temperatures. Great effort has been devoted to the oxidation behavior of ZrB2 based UHTCs and much progress has been obtained [9-16]. It has been shown that the introduction of SiC the ZrB2 ceramics results in significant improvement in oxidation resistance performance and mechanical property in comparison to ZrB2 alone, and ZrB2-SiC ceramic composites are currently considered the baseline UHTCs [17-20]. Numerous investigations have been conducted on the oxidation behavior of ZrB2-SiC ceramics, but these have mainly focused on static or flowing air studies under ambient pressures. The investigation to the oxidation behavior of the ceramic compos * Corresponding author. Tel./fax: +86 10 82316976. E-mail address: zhangy@buaa.edu.cn (Y. Zhang).  0010-938X/$ see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.004  ites under low oxygen partial pressure is needed in order to strengthen the understanding of the oxidation process of the materials in re-entry environments. Recently some researches have evaluated the oxidation resistance properties of the ceramic in simulated re-entry condition, and arc jet method was employed [21,22]. However, the oxygen partial pressure and chemical composition of the atmosphere can hardly be precisely controlled due to the rapid variation of the arc jet atmosphere, thus the effect of oxygen partial pressure on the oxidation process is still not clearly understood. Han and Fahrenholtz reported the oxidation behavior of ZrB2-SiC under low oxygen partial pressure, and their results indicated lower oxygen partial pressure is beneficial to the active oxidation of SiC. Hence the oxidation resistance performance is deteriorated compared to those oxidized in air [23,24]. However, both of researches by Han and Fahrenholtz focused on the effect of oxygen partial pressure on the evolution of the oxidation layer, and formation and evolution of SiC depleted layer was given detailed explanation. While the oxidation products had thought to be zirconia and borate silicate glass, thus the evolution of oxidation products has drawn less attention. As reported by Li [25] oxidation process of ZrB2-SiC ceramics is controlled by the diffusion of oxygen through oxidation layer and the latter is related to the formation and evolution of oxidation products during oxidation process. Therefore, a more comprehensive understanding of the formation and evolution of oxidation products during oxidation under low oxygen partial pressure is desired. The objective of this paper is to report formation and evolution of oxidation products for the ZrB2-SiC ceramics when oxidized under low oxygen partial pressure. The effects of oxygen partial pressure and oxidation temperature on the oxidation behavior of  \\x0c', '3298  D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  the ceramics are analyzed and the results obtained could provide a clearer insight into the oxidation process of the ceramics.  2. Experimental procedure  2.1. Materials processing  Commercially available ZrB2 and SiC powders with a reported purity of 99% and an average particle size of 2 lm were used. Ceramic composites containing of 80 vol.% ZrB2 and 20 vol.% SiC were prepared by hot-pressing method. The mixed powders were hotpressed in graphite dies lined with graphite foil and coated with BN. The mixtures were hot-pressed at 1900 °C for 30 min at a pressure of 30 MPa. The furnace was heated to the targeted temperature at an average rate of 30 °C/min after loading to the graphite dies, and argon was used as protective gas. The load was removed when the die temperature dropped below 1750 °C. The furnace was cooled to the room temperature when hot-pressing time was elapsed. Then the sintered specimens were put out from the graphite dies. Sample coupons in the size of 10 \\x02 10 \\x02 5 mm were cut from sintered materials by electrical discharge machining (EDM) method, and all the surface of the specimens were diamond polished to one finish. The coupons were ultrasonically cleaned successively in deionized water and ethanol before oxidation test.  and pure nitrogen was introduced as protective gas. Oxygen was introduced when reaching targeted temperature. After the expected oxidation time was arrived, the introduction of oxygen was stopped and the samples were cooled in nitrogen atmosphere. The oxygen partial pressure was set at 200 Pa during oxidation and the total pressure was fixed at 20 kPa. For comparison, oxidization in air of the ceramics was also conducted.  2.3. Characterization of oxidized specimens  Crystalline phases of oxidation products were determined by Xray diffraction (XRD, Dmax-2200, Rigaku, Tokyo, Japan) with Cu Ka radiation. The morphology and chemical composition of oxidation products on the surface and oxidation layer was detected by an Electronic Scanning Microscopy (S530, Hitachi, Tokyo, Japan) which equipped with an EDS detector (INCAINCAPentaFET-x3, Oxford, England). In addition, the micro-structure of oxidation products and oxidation layer was observed by an electrical probe micro-analyzer (EPMA, JXA8100, JEOL, Tokyo, Japan).  3. Results and discussion  3.1. Formation of zircon  2.2. Oxidation of ZrB2 - 20 vol.% SiC ceramics  The oxidation experiment was performed in a tube furnace with resistive molybdenum disilicide (MoSi2) heating elements capable of operating at 1800 °C. Oxygen (>99.99%) and nitrogen (>99.99%) were introduced at one end of the alumina tube and oxygen partial pressure can be adjusted by controlling the flow rate ratio between oxygen and nitrogen. Gas pressures in the alumina tube were measured and monitored by a digital pressure gauge. The samples were heated to a certain temperature of 1000, 1200, 1400, or 1600 °C at a ramp rate of approximately 5 °C/min,  For the materials oxidized under different temperatures for 30 min, the XRD patterns are collected and presented in Fig. 1, from which a notable phenomenon occurs while oxidation temperature arises to 1200 and 1400 °C, i.e., the emergence of zircon phase ZrSiO4. Additionally, the relative intensity of zircon is higher for the specimen oxidized under 1400 °C. Oxidation of silicon carbide can be omitted after oxidation at 1000 °C for 30 min, and the oxidation products are mainly composed of zirconia. Furthermore, most of the zirconium diboride is unaffected. After oxidation at 1600 °C, no oxidation products other than zirconia can be identified from the XRD pattern.  Fig. 1. XRD patterns of oxidized specimens, the symbols of B, Z, and S donate to zirconium diboride, monolithic zirconia and zircon respectively, and an enlarged XRD pattern of specimen oxidized at 1000 °C with 2h between 30° and 40° is offered.  \\x0c', 'D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  3299  between X-ray scattering coefficients for various phases. The XRD pattern of corundum phase Al2O3 is used for calibration and reference intensity ratio (RIR) method is employed to normalize the spectra obtained [26]. In addition, the zircon yield for the specimens oxidized in air was calculated in order to identify the effect of the oxygen partial pressure on the formation of zircon during oxidation process.  aZrSiO4 ¼  IZð2 0 0Þ IZð2 0 0Þ þ IT ð1 0 1Þ þ IMð1 1 1Þ þ IMð1 1 \\x161Þ  ð1Þ  As can be seen from the variation of zircon yield (shown in Fig. 2), both of oxidation temperature and oxygen partial pressure have obvious effect on the formation of zircon. The zircon yield reaches to its summit value when oxidation at 1400 °C under PO2 = 200 Pa, and the maximum value is about 0.595. Then it decreases with increment of oxidation temperature. On the other hand, the value of zircon yield for the specimens oxidized under low oxygen partial pressure is larger than those oxidized in air, and the zircon can be formed within a wider range of temperature when oxidation under low oxygen partial pressure. As shown in Fig. 2, the zircon yield reaches to 0.278 after oxidation at 1200 °C under a low oxygen partial pressure of 200 Pa, while no zircon can be detected from those oxidized in air at the same temperature. The variation trend of zircon yield indicates low oxygen partial pressure is beneficial to the formation of zircon, and the reason will be clarified in detailed in the following discussion. The elemental distribution of zirconium, silicon and oxygen on the surface of specimen oxidized at 1400 °C are shown in Fig. 3. As shown in the images, two different regions can be identified from the BSE image. The grey particles with regular columnar shape are larger where zirconium and silicon distributes homogeneously, and the white particles are smaller in size and mainly composed  Fig. 2. Variation of zircon yield during oxidation in air and under low oxygen partial pressure of PO2 = 200 Pa.  aZrSiO4  The zircon yield is calculated according to expression (1), and the qualitative variation trend of the zircon formation during oxidation process can be deduced from the results. In this expression, refers to the zircon yield, and the fraction of zircon in the crystal products is evaluated from the relative intensities, I, of the X-ray diffraction peaks for zircon (2 0 0) and tetragonal (1 0 1) and monoclinic (1 1 1, 1 1 \\x161) zirconias. No crystalline phase oxidation products other than zirconia and zircon can be identified from the XRD patterns, thus the zircon yield can be simplified according to the relation between zirconia and zircon. Herein, the value of zircon yield reflects the weight ratio of zircon among all crystalline oxidation products, i.e. zircon and zirconia. Furthermore, the XRD patterns should be dealt with due to differences  Fig. 3. Elemental analysis results to the specimen oxidized at 1400 °C.  \\x0c', '3300  D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  of zirconium. In addition, the atomic ratio between silicon and zirconium of each point are calculated according to quantitative analysis to the grey particles and white particles, and the results are shown in Fig. 3(d). The image shows that the atomic ratio between Si and Zr in the region of grey particles is near to 1, while no silicon can be detected from the whiter particles. All the elemental analysis results confirm our hypothesis that the zircon is formed during oxidation process, which is coincident with the XRD patterns. Fig. 4 shows the surface morphology of the ceramic composites after oxidation under a low oxygen partial pressure of 200 Pa. The surface of specimen oxidized at 1000 °C is covered with porous zirconia and unaffected SiC, and little silica glass can be observed, which can be confirmed from the EDS spectrum shown in Fig. 5. In addition, zircon particles are absent without participation of oxidation products of SiC, i.e. SiO or SiO2. Three phases of oxidation products can be identified from the surface of specimen oxidized at 1200 °C, as shown in Fig. 4(b). The white part donates to zirconia particles, the grey part refers to zircon crystals which have a mean size of 2 lm, and the black part is amorphous silica glass. Then the surface is almost covered by prism like zircon particles after oxidation at 1400 °C, and the mean size of the particles reaches to 10 lm (shown in Fig. 4(c)), which indicates increasing of oxidation temperature promotes the growth of zircon particles. Furthermore, the silica glass can be observed form the grain boundary of zircon. Fig. 4(d) shows the morphology of specimen oxidized at 1600 °C, the surface is covered by zirconia particles, and the silica glass distributes in the grain boundary of zirconia particles. The absence of zircon should be caused by the instability of zircon under this oxidation condition.  3.2. Formation mechanism of zircon  The probable formation routes of zircon are expressed in Eqs. (2) and (3), and the latter is related with the active oxidation of SiC, which is expressed as Eq. (4). The standard Gibbs energy of Eqs. (2) and (3) are calculated and the results are illustrated in Fig. 6. Herein the standard state means total pressure keeps at 1 atm. It is noted that the Gibbs energy of zircon produced by the gaseous-solid reaction is lower than the reaction (2) and it  Fig. 5. EDS spectrum of silica glass and unaffected SiC on the surface of specimen oxidized at 1000 °C for 30 min.  decreases with increasing of oxidation temperature, while the formation energy calculated from Eq. (3) increases obviously at elevated temperatures. The lower formation energy indicates that the second route is preferred for the formation of zircon, and the formation of gaseous SiO is necessary. According to the volatility diagram of ZrB2-SiC system calculated by Fahrenholtz [13], SiC exhibits preferential oxidation in the system of ZrB2-SiC, and lower oxygen partial pressure leads to the active oxidation of SiC under relatively lower temperature (as shown in Eq. (4)). In addition, the formation of protective silica film on the surface further lowers  Fig. 4. Surface morphology of specimens oxidized under PO2 = 200 Pa, and the oxidation temperatures are (a) 1000, (b) 1200, (c) 1400 and (d) 1600 °C.  \\x0c', 'D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  3301  Fig. 6. Standard Gibbs energy of two formation routes of zircon.  the oxygen partial pressure in the inner part, which aggravates the active oxidation of SiC. Hence formation of zircon by the second route is fulfilled with the participation of gaseous SiO.  ZrO2 ðcrÞ þ SiO2 ðlÞ ¼ ZrSiO4 ðcrÞ  ZrO2 ðcrÞ þ SiOðgÞ þ 1 2  O2 ðgÞ ¼ ZrSiO4 ðcrÞ  SiCðcrÞ þ O2 ðgÞ ¼ SiOðgÞ þ COðgÞ  ð2Þ  ð3Þ  ð4Þ  A schematic diagram of formation process of zircon during oxidation of ZrB2-SiC ceramics is demonstrated in Fig. 7, and the formation of zircon can be divided into two steps: (I) nucleation and (II) crystal growth. The first step is complete by the reaction (3), which is related to the active oxidation of SiC. The active oxidation of SiC produces gaseous SiO, then the gaseous SiO migrates to the underneath layer and reacts with zirconia, thus zircon micro crystalline is formed, as shown in Fig. 7(b). The oxygen diffuse to the underneath layer through protective silica rich film and reacts with the gaseous SiO, thus the thickness of the silica layer increases and the oxygen partial pressure under the underneath part is further lowered. As a result, the active oxidation of SiC is aggravated and more gaseous SiO will be emitted until the oxygen partial pressure is too small to provoke the active oxidation of SiC. The liquid silica layer promotes mass transferring of zircon micro crystalline and accelerates crystal growth of zircon, as shown in Fig. 7(c). Zircon is easily grown into prism like crystals according to its crystalline habits, as reported by Vavra [27]. To confirm the formation mechanism proposed in Fig. 7, the micro-structure evolution of cross section during oxidation process is analyzed, as shown in Fig. 8. A thin oxidation layer can be observed on the surface of the specimen oxidized under 1000 °C (Fig. 8(a)), which is composed of about two layers: (1) layer of amorphous borate glass, which is formed by oxidation of zirconium diboride; (2) porous zirconia layer. Detailed information about the porous zirconia layer can be obtained from an enlarged image offered in the right figure of Fig. 8(a). With oxidation temperature arises to 1200 °C, the amorphous boron oxide evaporates rapidly and disappears. The oxidation layer is composed of three layers: (1) layer of mixed crystals with thickness of about 2 lm, and the crystals on the surface can be divided into two phases, as shown in the surface image (Fig. 4(b)), thus zircon is loose and discontinuous. An enlarged image of this layer is shown in the right image of Fig. 8(b). (2) Mixed layer of zirconia particles and silica glass. (3) SiC depleted layer. After oxidation at 1400 °C, the oxidation layer becomes remarkably thinner (about 40 lm) and a layer of prism  Fig. 7. Schematic diagram of formation process of zircon (a) before oxidation; (b) reaction between gaseous SiO and zircon, and nucleated from silica glass; (c) growth of zircon crystalline into prism like crystals; and (d) zircon cannot be formed due to chemical instability at higher temperature.  like zircon with a thickness of about 10 lm can be observed, as shown in Fig. 8(c). Additionally, a layer of SiC depleted layer can be identified between the zircon layer and unaffected part, and this layer is also significantly thinner than that oxidized at 1200 °C. It should be noted that the oxidation process is accelerated at 1400 °C compared with that oxidized at 1200 °C, thus the surface of the specimen is covered with a layer of silica glass with lower viscosity, which aggravates active oxidation of SiC. The emission of SiO is related to the formation of zircon, as indicated in Eq. (2). The formation of the dense and continuous silica glass layer with zircon crystals dispersants lowers diffusion rate of oxygen according to the investigation result of Cherniak [28,29], and the oxygen partial pressure in the inner part of the ceramic is further lowered until the oxygen partial pressure is too low for the active oxidation of SiC. Hence the SiC depleted layer is thinner than that oxidized under 1200 °C. The zircon microcrystalline grow into prism like particles with the help of silica glass with extension of holding time, as shown in an enlarged image offered in Fig. 8(c),  \\x0c', '3302  D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  Fig. 8.  Images of cross section of ZrB2-20 vol.% SiC ceramics oxidized in air for 30 min, oxidation layers are marked respectively. (a) 1000, (b) 1200, (c) 1400 and (d) 1600 °C.  and a dense and continuous zircon layer is formed finally. After oxidation at 1600 °C for 30 min, zircon cannot be obtained due to its chemical instability, and the oxidation layer is composed of about three layers, (1) columnar zirconia layer with silica glass dispersed in the grain boundary. (2) Mixed layer of zirconia and silica glass, with a thickness of 30 lm; (3) SiC depleted layer, with a thickness of about 50 lm.  3.3. Oxidation resistance performance of zircon  The oxygen diffusion rate in zircon is far less than that of zirconia, as reported by researchers [27,28], then oxidation resistance performance of zircon should be better than that of zirconia. In this section, we evaluate the oxidation resistance performance of zircon according to the variation of the thickness of oxidation layer. The variation of thicknesses for the oxidation layer, SiC depleted layer and zircon layer are measured in accordance to Fig. 8 and the results are shown in Fig. 9. As noted by the image, the variation of the thickness of oxidation layer is related to the formation of zircon. The oxidation layer is merely 10 lm after oxidation at 1000 °C for 30 min, which is mainly composed of porous zirconia and discontinuous borate glass. With the oxidation temperature rises to 1200 °C, the thickness of the oxidation layer rises to  Fig. 9. Variation of thickness for the oxidation layer, SiC depleted layer and zircon layer after oxidation under various temperatures.  72 lm rapidly. The sharply increasing of oxidation layer is related to the discontinuous zircon on the surface, and oxygen can diffusion into the inner part of the ceramic from the site of zirconia particles. A layer of zircon with a thickness of about 10 lm can be observed for the specimen oxidized at 1400 °C, and this layer is  \\x0c', 'D. Gao et al. / Corrosion Science 52 (2010) 3297-3303  3303  dense and continuous, as shown in Fig. 4(c), which improves the oxidation resistance performance of the ceramic. The thickness of SiC depleted layer has a thickness of 20 lm, the thickness of the oxidation layer decreases to 40 lm, which is significantly thinner than that oxidized at 1200 °C. With oxidation temperature increases to 1600 °C, the thickness of oxidation layer increases rapidly to 160 lm due to instability of zircon at this temperature. The oxidation layer consists of a layer of columnar zirconia with a thickness of about 80 lm, a layer of zirconia and silica glass coexistence layer, which has a thickness of about 30 lm, and the SiC depleted layer with a thickness of about 50 lm. Interestingly, we found that the variation of thickness of oxidation layer for the specimen oxidized at 1000, 1200 and 1600 °C, i.e., the oxidized specimens without a dense and continuous zircon layer, follows a linear relation with oxidation temperature. And the formation of zircon improves oxidation resistance performance of the ceramic obviously, as shown in Fig. 9. This can be attributed to the formation of dense and continuous zircon layer, which lowers diffusion rate of oxygen significantly compared to that of zirconia. The variation trend of oxidation layer confirms that formation of zircon is beneficial to improving the oxidation resistance performance of ZrB2-SiC ceramics.  4. Conclusion  Zirconium diboride ceramics containing 20 vol.% SiC are oxidized under a low oxygen partial pressure of PO2 = 200 Pa. Formation and evolution of zircon during oxidation process is given thoroughly investigation, which has been mentioned scarcely in previous research. The experimental results confirm the formation of zircon for the ceramics oxidized under 1200 and 1400 °C, and the formation process can be divided into two steps (1) nucleation from silica glass, which is completed by reaction between gaseous SiO and zirconia; (2) crystal growth into prism like particles, which is fulfilled in the liquid silica glass. The effect of oxidation temperature and oxygen partial pressure on the formation and evolution of zircon is analyzed, and the experimental results indicate that the chemical stability of zircon decreases with increasing of oxidation temperature, and lower oxygen partial pressure is beneficial to the formation of zircon due to the active oxidation of SiC. A dense and continuous layer of zircon is obtained for the specimen oxidized at 1400 °C under PO2 = 200 Pa, which significantly improves the oxidation resistance performance of the ceramic.  References  [1] K. Bundschuh, M. Schutze, Materials for temperatures above 1500 °C in oxidizing atmospheres. Part I: basic considerations on materials selection, Materials and Corrosion 52 (2001) 204-212. [2] Frederic Monteverde, Raffaele Savino, Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions, Journal of the European Ceramic Society 27 (2007) 4797-4805. [3] Frederic Monteverde, Luigi Scatteia, Resistance to thermal shock and to oxidation of metal diborides-SiC ceramics for aerospace application, Journal of the American Ceramic Society 90 (2007) 1130-1138.  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Hilmas, Oxidation of zirconium diboride-silicon carbide at 1500 °C at a low partial pressure of oxygen, Journal of the American Ceramic Society 89 (2006) 3240-3245. Ju Li, Thomas J. Lenosky, et al., Thermochemical and mechanical stabilities of the oxide scale of ZrB2-SiC and oxygen transport mechanisms, Journal of the American Ceramic Society 91 (2008) 1475-1480. [26] Linhong Guo, Hui Li, Phase transformations and structure characterization of calcium polyphosphate during sintering process, Journal of Materials Science 39 (2004) 7041-7047. [27] Gerhard Vavra, On the kinematics of zircon growth and its petrogenetic significance: a cathodoluminescence study, Contrib Mineral Petrol 106 (1990) 90-99. [28] Daniele J. Cherniak, E. Bruce Watson, Diffusion Mineralogy and Geochemistry 51 (2003) 113-143. [29] F.J. Keneshea, D.L. Douglass, The diffusion of oxygen in zirconia as a function of oxygen pressure, Oxidation of Metals 3 (1971) 1-14.  zircon,  Reviews  in  [18]  in  \\x0c']"
},{
  "_id": 177,
  "PDF": "Oxidation of ZrB2 and its composites - a review.pdf",
  "Text": "['R E V I E W  Oxidation of ZrB2 and its composites: a review  Ryo Inoue1,*  , Yutaro Arai2, Yuki Kubota3, Yasuo Kogo1, and Ken Goto1,4  1 Department of Materials Science and Technology, Tokyo University of Science, 6-3-1, Niijyuku, Katsushika-ku, Tokyo 125-8585,  Japan  2 Katayanagi Advanced Research Laboratories, Tokyo University of Technology, 1404-1 Katakuramachi, Hachioji, Tokyo 192-0982,  Japan  3 Structures and Advanced Composite Research Unit, Japan Aerospace Exploration Agency (JAXA), 2-12-1, Jindaizi-higashimachi,  Chofu-shi, Tokyo 182-0021, Japan  4 Institute of Space and Astronautical Science(ISAS), Japan Aerospace Exploration Agency (JAXA), 3-1-1, Yoshinodai, Chuo,  Sagamihara, Kanagawa 252-5210, Japan  Received: 10 March 2018  Accepted: 15 June 2018  Published online:  25 June 2018  Ó Springer  Science+Business  Media, LLC, part of Springer  Nature 2018  ABSTRACT  The  oxidation behavior  and oxidation mechanisms  of monolithic ZrB2  and  particulate-ZrB2 matrix composites were reviewed. Dispersion of SiC particles  into ZrB2 was  found to be an effective way to prevent  extensive oxidation.  However,  the formation of a SiC-depleted layer can become a critical problem  because it can lead to spallation and delamination of the protective surface layer.  The addition of ZrC in conjunction with rapid heating to temperatures higher than 2000 °C effectively reduced the porosity of  the SiC-depleted layer. The  formation of a dense surface layer was attributed to large volumetric expansion  during  the  conversion  from ZrC to ZrO2. The  effect  of  the ZrC addition  depended on the  temperature, heating  rate,  and composition. This  review  showed that material design for specific applications is required for high-tem perature applications to maximize the oxidation resistance of ZSZ composites.  Application and oxidation of monolithic  zirconium diborides  Diborides, carbides, and nitrides of refractory metals  (Zr, Hf, Ta, etc.) and their  composites, which have very high melting points ([ 3000 °C [1-24] or [ 2500 °C [16, 25-29]), are categorized as ultra-high temperature  ceramics  (UHTCs). UHTCs have been  used as heat-resistant materials  in extremely harsh  environments,  such as  the nose  cone  and leading  edge of hypersonic vehicles [30],  thermal protection  systems  (TPS)  for  aerospace  re-entry  vehicles  [15, 31-37], propulsion systems [31], furnace elements  [32],  refractory crucibles  [32], and plasma-arc  elec trodes [33]. Figure 1 shows the relationship between melting point (Tm) and density (q) for typical UHTCs  [34]. Research  on  and  development  of  transition metal diborides,  carbides, nitrides, oxides, and sili cides  began  in  the  1940 s  and 1950 s.  Pioneering  research by Norton et  al.  [35]  and Samsnov et  al.  [36, 37] led to the determination of physical constants,  such as the lattice constants and the densities of such  Address correspondence to E-mail:  inoue.ryo@rs.tus.ac.jp  https://doi.org/10.1007/s10853-018-2601-0  J Mater Sci  (2018) 53:14885-14906  Review  \\x0c', 'materials, and examination of  fundamental proper ties,  such as  the  thermal  conductivities of various  borides. Many studies on UHTCs were carried out in  the 1960 s and 1970 s [38-51]. These works  focused  on refractory metal diborides, such as ZrB2 and HfB2,  and their composites.  Among UHTCs, ZrB2 has quite a high melting point ([ 3000 °C [7, 9, 11, 24, 52-83]) and a relatively  low density. Thus, ZrB2 and its composites have been  considered to be among the most attractive candidate  materials  for  heat-resistant  aerospace  components.  However, oxidation of ZrB2 is a critical problem for  the above-mentioned high-temperature applications.  ZrB2  is  oxidized when  exposed  to  an  oxidizing  atmosphere,  in  accordance  with  the  following  reaction: Þ þ 5=2O2 g  ZrB2  sð  \\x00  \\x01  ! ZrO2  sð  Þ þ B2O3  lð Þ  ð1Þ  The vapor pressure of B2O3  is  among the highest  vapor pressures for oxides [84]. B2O3 vaporizes above * 1100 °C at atmospheric pressure, as follows: lð Þ ! B2O3 g  B2O3  \\x00  \\x01  ð2Þ  The oxidation behavior of ZrB2 strongly depends  on temperature. Various researchers have examined  the oxidation behavior of monolithic ZrB2 in a wide to * 1500 °C).  range  of  temperatures  (from 500  Kuriakose et al. [38] oxidized ZrB2 at temperatures in the range of 945-1256 °C. Their research showed that  the oxidation rate of ZrB2 obeyed a parabolic law in  this temperature range. The rate of parabolic oxida tion of ZrB2 (hereinafter denoted by found to be 0.1019 (mg/cm2)2/min at 945 °C, and it  kp;ZrB2 ) was  increased to 0.5625 (mg/cm2)2/min at 1256 °C under a constant oxygen partial pressure, PO2 , of * 105 Pa  (740 ± 5 mm Hg).  Voitovich et al. [85] heated ZrB2 at temperatures in the range of 300-1200 °C for 6 h in air and calculated  the isobaric-isothermal potential  (Gibbs free energy)  for  the reaction of ZrB2 with air. Oxidation of ZrB2  was observed to begin at temperatures in the range of 500-550 °C, with no microstructural changes observed at temperatures up to 500 °C. A smooth and  golden-bluish film, composed of B2O3, covered the specimen surface after heating at 550 °C. At higher temperatures, in the range of 650-780 °C, oxidation of  ZrB2 still obeyed a parabolic law. First, ZrO2 with a  B2O3  layer  formed on the  surface, and then liquid  B2O3 was found to exist  in a porous ZrO2 skeleton.  From the  perspective  of  oxidation  resistance,  the  temperature  at which B2O3  starts  to  evaporate  is  important  for understanding the oxidation mecha nisms of ZrB2 and its composites. At temperatures in the range of 780-950 °C,  the trend exhibited by the  oxidation rate changed from parabolic to cubic. The  oxidation rate tended to be linear at temperatures above 1200 °C because of microstructural changes on  the  surface,  i.e.,  cracking  and  peeling  of  surface  scales. Figure 2 presents crack evolution from ZrO2 layer after oxidation at * 1532 °C for 30 min (Fig. 2a)  and the changes in morphology at the boride/oxide for ZrB2 after oxidation at * 1500 °C for  interface  30 min  (Fig. 2b)  [41,  86].  Figure 2a  shows  crack  nucleation in the surface scales, with oxygen pene trating  the  unoxidized  region  through  pores  and  cracks  in the defective  scales. Figure 2b shows  the  formation of porous ZrO2 on ZrB2. As porous ZrO2  does  not  act  as  a  barrier  for  oxygen  diffusion,  Figure 1 Relationship  between melting  point  and  density  for  various UHTCs [34].  500μm  (a)  20μm  ZrO2  ZrB2  (b)  ZrO2  ZrB2  Figure 2 a Crack evolution from ZrO2 * 1532 °C for 30 min and b changes  layer after oxidation at  in morphologies  at  the  boride/oxide interface of polycrystalline ZrB2 after oxidation at * 1500 °C for 30 min [41, 86].  14886  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'oxidation of ZrB2  is accelerated not only by the dis appearance of  the B2O3  layer but also by the occur rence of crack penetration in this temperature region.  Mattuck [41] oxidized ZrB2 in the temperature range of 927-1545 °C and observed that  the rate of oxygen  consumption during oxidation increased rapidly above * 1200 °C, probably because B2O3 vaporized  as soon as an oxide layer (mainly composed of B2O3  and ZrO2) was formed. Tripp and Graham [48] oxi dized ZrB2 at temperatures in the range of 800- 1400 °C at a PO2 of 3.3 9 104 Pa (250 mm Hg) using a  thermogravimetric  (TG)  furnace.  They measured  both the vaporization rate of B2O3 and the oxidation  rate  of  ZrB2.  Figure 3  exhibits  the  relationship  between the oxidation rate of ZrB2  (kp;ZrB2 ) and the  vaporization rate of B2O3  (kvap )  [38, 41, 48]. Figure 3  clearly shows that kp;ZrB2  increases with the increase  in kvap . Based on their measurements, they developed  the following expression for the weight gain of ZrB2 (DwZrB2 ) during oxidation: ¼ 17:65 exp \\x00 2:5 \\x02 104 cal þ 9:15 \\x02 104 exp \\x00 4:7 \\x02 104 cal  DwZrB2 mg/cm2  \\x10  \\x11  ð  Þ  2RT  \\x12  \\x13  t1=2  ð  Þ  RT  \\x12  \\x13  t  ð3Þ  The rate of  total oxygen consumption during oxi dation of ZrB2 (CZrB2 ) can be expressed as follows:  CZrB2 mol/cm2  \\x10  \\x11  ¼ 5:41  \\x004 exp \\x00 2:5 \\x02 104 cal \\x02 10 þ 2:86 exp \\x00 4:7 \\x02 104 cal  ð  Þ  2RT  \\x12  \\x13  t1=2  ð  Þ  RT  \\x12  \\x13  t  ð4Þ  where t, T, and R are the oxidation time, temperature,  and gas constant, respectively. perature range of 800-1500 °C,  In fact,  in the  tem the oxidation rate is 1100 °C,  governed  by  a  paralinear  law above  as  shown  in Eq. 1,  because  the  evaporation  of B2O3  affects  the oxidation rate of ZrB2. Figure 4 summa rizes  the weight gain of ZrB2 during heating. The  oxidation behavior of ZrB2 is strongly dependent on  experimental conditions such as experimental setup,  specimen shape, gas flow rate, and specimen place ment method. In particular,  it appears that diffusionabove * 1200 °C  limited oxidation of ZrB2 occurs  [38, 41, 48].  In addition, PO2 was also found to have a strong  effect on the oxidation rate in the system. Figure 5  illustrates the increase in kp;ZrB2 with increasing PO2 [38]. The parabolic oxidation rate for ZrB2 at 945 °C is * 0.025 (mg/cm2)2/min when * 1.4 9 104 Pa (* 102 mm Hg), and it increases up to * 0.2368 (mg/cm2)2/min at * 105 Pa (* 740 mm  PO2  is  less  than  Hg). At a given temperature, the weight gain of ZrB2  during heating can vary because the oxidation rate of  ZrB2 depends on both the temperature and PO2 .  These  experimental  results  showed that  the oxi dation  of ZrB2  can  be  divided  into  two  phases:  Figure 3 Relationship between oxidation rate of ZrB2  (kp;ZrB2 )  and vaporization rate of B2O3 (kvap )  [38, 48].  Figure 4 Weight gain of ZrB2 during heating as a function of  heating time,  t [45, 48, 100].  J Mater Sci  (2018) 53:14885-14906  14887  \\x0c', 'below 1100 °C and (I) (II) above 1100 °C, B2O3 acts as a barrier  1100 °C.  Below  to oxygen diffusion  from the surface toward the unoxidized region. B2O3 starts to evaporate at temperatures above 865 °C, and  oxidation  of  ZrB2  proceeds  by  oxygen  diffusion  through the ZrO2 layer in the temperature range of 950-1100 °C. Above 1100 °C, significant evaporation  of B2O3 occurs, and B2O3 does not  form protective  oxide scales on the surface. Thus, simultaneous oxi dation and morphological changes occur. In terms of  oxidation kinetics, parabolic oxidation behavior cea ses to occur  (or  rather,  the linear component of  the  paralinear oxidation behavior increases drastically) in  this temperature region.  Based on  experimental  results,  Fahrenholtz  [86]  calculated the vapor pressure of the oxide evolved by  oxidation of ZrB2 and presented volatility diagrams  for ZrB2. Figure 6 shows volatility diagrams for ZrB2 at 727, 1527, and 2227 °C. At 727 °C,  the dominant  gaseous species is BO2 (g), with a vapor pressure on the order of 10-6 Pa. B2O3 is maintained as a liquid on  the surface of oxidized ZrB2, where it acts as a barrier  to oxygen diffusion. However,  the vapor pressure of  B2O3  increases  rapidly  at  higher  temperatures.  Fahrenholtz concluded that the oxidation rate of ZrB2 temperatures above * 1370 °C  increases  rapidly at  owing to evaporation of the B2O3 layer. A B2O3 layer is not formed on the oxidized ZrB2 surface at 1527 °C  because the vapor pressure of B2O3 becomes quite high (* 344 Pa) at  this temperature.  Parthasarathy et al.  [87]  calculated the oxidation  rate,  oxidized layer  thickness,  and ZrB2  recession  using  a model  constructed  from the  results  of  a  thermodynamic  analysis.  They  assumed  that  the  oxidation of ZrB2  resulted in liquid B2O3 and ZrO2  existing on an oxidized ZrB2  surface. Figure 7 pre sents a comparison between the calculated oxidation  rate and experimental data as a function of  temper ature. The oxidation rates calculated as a function of  thermodynamic parameters  such as PO2 ,  the partial  pressure  of B2O3,  and the  oxygen permeability of  B2O3 are in good agreement with the experimental  data. This  type of  theoretical  analysis  explains  the  Figure 5 Relationship between oxidation rate of ZrB2  (kp;ZrB2 )  and oxygen partial pressure (PO2 )  [38].  −30.0 −40  B(g)  BO(g)  B2O2(g)  B2O3(g)  BO2(g)  ZrO2 (cr) + B2O3(l)  ZrB2 (cr)  727°C  1527°C  2227°C  −20.0  −10.0  0.0  10.0  g o L  p  (  x  )  −30  −20  −10  0  10  Log pO2  Figure 6 Volatility diagrams for ZrB2 at 727, 1527, and 2227 °C  [86].  Figure 7 Comparison  between  calculated  oxidation  rate  and  experimental data as a function of  temperature [87].  14888  J Mater Sci  (2018) 53:14885-14906    \\x0c', 'relationship between the oxidation rate of ZrB2 and  temperature well.  In addition, polycrystalline ZrB2 is difficult  to fab ricate by sintering, and residual porosity also affects  the oxidation behavior. For example, Kuzenkova and  Kislyi  [39] oxidized ZrB2 with porosities of 8% and 1000 °C for  15% at  10 h. They confirmed that  the  porosity of ZrB2 affected the extent of weight gain  during oxidation, as the weight of ZrB2 with a porosity of 8% increased by * 5 mg/cm2 after 10 h at 1000 °C, whereas that of ZrB2 with a porosity of 15% increased by * 12 mg/cm2. The  relationship  between sintering method and relative density is well  summarized by Guo et al. [25] because the porosity of  ZrB2  affects  the weight  gain during  oxidation,  as  mentioned above. This  study showed that various  fabrication routes are available that can address the  process cost and the applicability to large and com plicated shapes, as well as improve the fundamental  properties,  including  strength,  toughness, weight  gain during oxidation, and creep resistance, at ele vated temperatures.  Thus, oxidation of monolithic ZrB2 is dominated by  various factors, such as temperature, oxygen partial  pressure,  and  morphology.  In  particular,  it  is  important  to form a protective surface layer on an  unreacted region to prevent degradation of  this pro tective  surface  layer  through cracking, peeling,  or  delamination.  In  this paper,  the  current  status  of  attempts  to prevent oxidation of ZrB2, as described  above,  is reviewed based on previous  research and  our recent studies.  Addition of a secondary phase to improve  the oxidation behavior of ZrB2  The addition of a secondary phase to polycrystalline  ZrB2 has been attempted for more than two decades  in order to withstand oxidation at temperatures above 1200 °C. In this section, we focus on the effect  of  the second-phase particles on the oxidation resis tance  of ZrB2. From the perspective of  improving  oxidation, materials  such  as  SiC,  transition-metal  disilicides (MeSi2), Si3N4, and Zr-based UHTCs (e.g.,  ZrC and ZrN)  [59, 62, 88-90] have been added to  ZrB2. Among these additives, previous studies have  mainly focused on the effects of SiC and MeSi2. The  oxidation rates of ZrB2-SiC and ZrB2-MeSi2 decrease  owing  to  the formation of SiO2 scales temperatures above 1200 °C: oxidation at SiC sð Þ þ 3=2O2 g ! SiO2 lð Þ þ CO g  by passive  \\x00  \\x01  \\x00  \\x01  ð5Þ  xMeSi2  sð  Þ þ 2x þ 1=2y  \\x00  \\x01  O2 g  \\x00  \\x01  ! 2xSiO2  lð Þ þ MexOy  ð6Þ  The oxygen permeability of vitreous SiO2 is much  lower than that of B2O3, and it is stable at tures above 1600 °C. The mechanical,  tempera thermal, and  electrical properties  and the oxidation behavior of  ZrB2-SiC (hereinafter  denoted  by ZS)  and ZrB2-  MeSi2  composites  have  also  been  studied  [13, 54, 83, 91, 92]. These composites have a critical  problem in terms of their oxidation behavior because temperatures above 1500 °C, SiC and SiO2 cause  at  the evolution of gaseous SiO under  the SiO2  scales  [14], according to the following reaction: SiC sð Þ þ O2 g ! SiO g þ CO g  \\x00  \\x01  \\x00  \\x01  \\x00  \\x01  ð7Þ  SiO2  lð Þ ! SiO g  \\x00  \\x01  þ 1=2O2 g  \\x00  \\x01  ð8Þ  Reaction (7) is called active (preferential) oxidation of  SiC. A decrease in the oxygen partial pressure or an  increase  in  the  temperature  of  the  system causes  these  reactions  to  occur,  and SiO2  scales  are  not  formed or are decomposed.  -10  -20 -30  -20  0  10  -10  0  10  Oxygen partial pressure, PO 2, Log(Pa)  V  o p a  r  p  r  e  s s  u  r  e  ,  P  v  p a  ,  g o L  (  P  a  )  ZrB2 (s) + SiC (s)  ZrO2 (s)  + B2O3(l) + SiO2 (l)  SiO (g)  B2O3(g)  BO2(g)  ZrB2 (s) + SiO2 (l)  Figure 8 Volatility diagram for ZS composites at 1500 °C [14].  J Mater Sci  (2018) 53:14885-14906  14889        \\x0c', 'Addition of SiC  Figure 8 shows a volatility diagram for ZS composites at 1500 °C [14]. Active oxidation of SiC depends on PO2 and the temperature of the system. At 1500 °C,  SiO is formed by active oxidation of SiC when PO2 falls below 1.8 9 10-11 Pa. Even in an air atmosphere (PO2 = 2.0 9 104 Pa), SiO is formed at above * 1800 °C [14]. After  temperatures  formation of an oxide  layer  containing SiO2  (SiO2 ? ZrO2), PO2 under  the  oxide  layer decreases  because  of  the  low oxygen  diffusivity in SiO2 [84, 93].  Fahrenholtz [14] showed that preferential  (active)  oxidation of SiC proceeds under the oxide layer and  that a porous ZrB2 (and ZrO2) layer is formed under a  SiO2-based oxide layer after oxidation of ZS composites with 30 vol% SiC at 1500 °C for 30 min. Fig ure 9 shows cross sections of oxidized ZS composites (a) 1500 °C (30 vol% SiC), (b) 1800 °C (20 and 80 vol% SiC), (c) 1900 °C (30 vol% SiC), (d) 2000 °C (20 vol% SiC), and (e) 2200 °C (20 vol% SiC)  at  [14, 94-97].  As shown in Fig. 9a, a porous layer composed of ZrB2  (and ZrO2) develops between the SiO2 layer and the  unoxidized region, which is  referred to as a ‘‘SiC depleted layer.’’ Owing to its porous  structure,  the  existence of a SiC-depleted layer under SiO2-based  oxide  scales  causes  spallation and delamination of  the oxide scales formed during oxidation.  Williams et al. [94] fabricated ZS composites with a  wide range of SiC contents: SiC:ZrB2  (vol%) = 80:20  (S80Z20),  65:35  (S65Z35),  50:50  (S50Z50),  35:65  (S35Z65),  and 20:80  (S20Z80). They  evaluated the  effect of SiC content on the oxidation behavior of ZS  composites. Oxidation of these ZS composites at 1800 °C for 20 min revealed that a SiO2 ? ZrO2 layer  is formed on all ZS composites because SiO2 (l) on the  surface evaporates and resulting in ZrO2 under  the  SiO2  layer being exposed in this temperature range.  The thickness of SiC-depleted layer is 110 ± 10, 160 ±  15, and 170 ± 10 lm for S20Z80, S35Z65, and S50Z50,  respectively.  In contrast, a SiC-depleted layer  is not  formed for S80Z20 and S65Z35 because SiO2 (l) is evaporated from the surface (SiO2 (l) ? SiO (g) ? 1/  2O2  (g)), but  the oxygen partial pressure below the  scale is not low enough to form a SiC-depleted layer.  Although the oxidation behavior of ZS composites depends on the composition, at * 1800 °C, ZS com posites form SiO2 or SiO2 ? ZrO2 oxide layers. However, holding times of[ 20 min at 1800 °C result  in the disappearance of SiO2  from the surface of ZS  composites.  Han et al.  [95] oxidized ZS composites (SiC = 10, 1900 °C for  20,  and 30 vol%)  at  1 h in air.  It was  revealed that a SiO2-rich layer with a thickness of * 50 lm and a ZrO2-rich layer with a thickness of * 100 lm were  formed on  the ZS  composites  as  oxide layers and a SiC-depleted layer with a thickness of * 800 lm was also formed after oxidation  (Fig. 9c)  Subsequently,  cracks  and spallation were  observed in the SiC-depleted layer. Above * 1800 °C, the growth rate of the SiC-depleted layer  increased  because  active  oxidation  of  SiC  was  accelerated.  Inoue et al. [96] oxidized ZS composites with 10, 20, and 30 vol% SiC above 2000 °C for 5-10 s by con ductive heating. A ZrO2  layer was observed as an  oxide layer on the ZS composites, but a SiO2  layer,  which was formed 2000 °C, was not A SiC-depleted layer with a thickness of 10-30 lm  on  the ZS  composites  below  formed after  oxidation (Fig. 9d).  was formed after oxidation, even for time (5-10 s). Above 2000 °C, a SiO2  such a  short  layer was not  formed on the surface of the ZS composites owing to  the decomposition of SiO2 (l).  Han et  al.  [97] oxidized ZS 2200 °C using  composites with 20  vol% SiC at  an oxyacetylene  torch.  They reported the formation of a porous ZrO2 layer  and a recrystallized ZrO2 layer as oxide layers, with a  void layer and a SiC-depleted layer formed beneath  the ZrO2 layer (Fig. 9e). These findings show that it is  quite difficult  to apply ZS composites  to materials  exposed  to  oxidizing atmospheres, above * 2000 °C,  especially  at  temperatures  because  of  active  oxidation of SiC. The thickness of the oxidized layer,  tox, which is defined as the sum of the thicknesses of  all  the surface oxide scales  (SiO2  layer, ZrO2  layer,  and/or  SiO2 ? ZrO2  layer)  and  the  SiC-depleted  layer,  is summarized in Fig. 10. The thickness of the oxide layer is * 500 lm after oxidation for 50 h at * 1450 °C, * 800 lm after oxidation for 60 min at 1900 °C, and * 450 lm after oxidation for 10 min at 2200 °C. It is evident that drastic oxidation of ZS composites occurs at * 2000 °C and that  this phe nomenon limits the use of ZS composites at atures above 2000 °C.  temper 14890  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'Zr  O  Si  SEM  100μm  S20Z80 (80ZrB2-20SiC (vol%))  SiO2 layer  Un-oxidized  SiO2 + ZrO 2 layer  SiC-depleted layer  Un-oxidized  J Mater Sci  (2018) 53:14885-14906  Figure 9  SEM images and  elemental mapping of cross  sections of oxidized ZS composites: a 1500 °C (30 vol% SiC), b 1800 °C (20 and 80 vol% SiC), c 1900 °C (30 vol% SiC), d 2000 °C (20 vol% SiC), and e 2200 °C (20  vol% SiC,  torch test)  [14, 90-93].  (a)  SEM  (b)  14891  SiO2 layer  SiC-depleted layer  Un-oxidized  O  Si  10μm  S80Z20 (20ZrB2-80SiC (vol%))  Zr  (c)  (d)  BSE  (e)  O  Si  SEM  200μm  SiO2 (rich) layer  ZrO2 (rich) layer  SiC-depleted layer  SiC-depleted layer  300μm  Un-oxidized  Crack/spallation  300μm  Zr  O  Si  C  ZrO2 layer  SiC-depleted layer  Un-oxidized  ZrO2 layer  Recrystallized ZrO2 layer  Void layer (SiC-depleted layer)  SiC-depleted layer  50μm  Un-oxidized  \\x0c', 'Addition of metal disilicides  The effectiveness of MeSi2 addition to improve the  oxidation behavior of ZrB2 has also been examined  since the 1960 s. Kuzenkova and Kislyi [39] fabricated  (ZrB1.9)23-MoSi1.1 and (ZrB1.7)13-MoSi1.2 by sintering  of ZrB2 with 5 and 10% MoSi2  (porosity 5-9%) and  evaluated  their weight gain during oxidation temperatures up to 1600 °C in air. They weight gain of * 20 mg/cm2 for (ZrB1.9)23-MoSi1.1 1400 °C in  at  found a  and  (ZrB1.7)13-MoSi1.2  at  air  over  9 h,  which is half the weight gain observed for monolithic  ZrB2 under the same conditions. Thus, the addition of  MoSi2 to polycrystalline ZrB2 was concluded to improve the oxidation resistance in air up to 1600 °C.  Lavrenko  et  al.  [98]  also  fabricated  ZrB2-ZrSi2  binary composites with ZrSi2 contents of 5, 15, and 50  wt% (6, 18, and 55 vol% ZrSi2). They reported that the  addition of 6 and 18 vol% of ZrSi2 to ZrB2 improved 1000 °C because  the  oxidation  behavior  above  a  ZrSiO4  (zircon)  layer  is  formed on the ZrB2-ZrSi2  surface, which prevents further oxidation.  The effect of adding various disilicides to ZrB2  to  improve the oxidation behavior has been well docu mented by Sciti and co-workers [89]. Based on their  studies on the oxidation of disilicides containing ZrB2  [99], to clarify the effect of disilicide addition to ZrB2  on the oxidation behavior of ZrB2-MeSi2,  they pre pared four different ZrB2-15 vol% MeSi2 composites  (Me = Zr, Mo, Ta, or W) and these samples at 1200, 1350, 1500, 1650, and 1800 °C for 15 min. Figure 11  shows  the oxidation behavior of  the ZrB2-15 vol%  MeSi2 samples at 1200, 1350, 1500, and 1650 °C, and  the thickness of  the oxidized layer as a function of  temperature is summarized in Fig. 12 [89].  For ZrB2-ZrSi2, the thickness of the oxidized layer is * 30 lm, and the oxidized surface is covered with  B2O3 and ZrO2 contained in a SiO2 1200 °C because  layer after oxi dation  at  the  oxidation  of ZrSi2  produces SiO2, as follows: Þ þ 3O2 g ! ZrO2  ZrSi2  sð  \\x00  \\x01  sð  Þ þ 2SiO2  l; s  ð  Þ  ð9Þ  Oxides formed on the surface are not affected by oxidation, even at 1650 °C, but crystal growth of ZrO2  in the SiO2 layer is accelerated by increasing the oxidation temperature. After oxidation at 1650 °C, a  columnar-like ZrO2 phase is formed in the oxidized  layer and gaseous products (B2O3, SiO) are evacuated  thorough the SiO2 layer with columnar-like ZrO2.  For ZrB2-MoSi2, SiO2 layer with ZrO2 and B2O3 is  formed on the surface, and the formation of MoB is  also observed after oxidation at 1650 °C.  temperatures up to  The  following  reactions  are  expected  to  occur during oxidation: Þ þ 7=2O2 g ! Mo3 Si5  5MoSi2  sð  \\x00  \\x01  sð  Þ þ 7SiO2  l; s  ð  Þ  ð11Þ  Þ þ B2O3 lð Þ þ 5=2O2 g 2MoSi2 ! 4SiO2 l; s Þ þ 2MoB sð  sð  \\x00  \\x01  ð  Þ  ð12Þ  After oxidation at 1650 °C, the oxidized layer consists  of SiO2 with a ZrO2 layer and a ZrB2-SiO2-MoB layer.  The thickness of the oxidized layer after oxidation at 1650 °C is * 50 lm, which  is  comparable  to  the  thickness of  the oxidized layer formed on ZS com1800 °C, MoB  posites. After  oxidation  at  is  not  observed on the oxidized surface because oxidation  of MoB  occurs  in  accordance with  the  following  reaction: 2MoB þ 9=2O2 g  \\x00  \\x01  ! 2MoO3  sð  Þ þ B2O3  lð Þ  ð13Þ  MoO3 was found to form between SiO2 and ZrO2 on  the external surface of the oxide layer after oxidation at 1800 °C, and after oxidation at 1800 °C, ness of the oxidized layer was * 180 lm. For ZrB2-TaSi2, after oxidation at 1200 and 1350 °C,  the thick the oxidized surface is covered with tiny ZrO2 par ticles contained in a SiO2  layer.  In this case, SiO2  is  supplied by the following reaction: Þ þ 13=2O2 g ! Ta2O5 Þ þ 4SiO2  TaSi2  sð  \\x00  \\x01  sð  l; s  ð  Þ  ð14Þ  In addition,  the  following complex reactions occur  Figure 10 Thickness of oxidized layer on ZS composites  as  a  function  of  temperature, T. The  times  in  this  graph  show the  holding time at each temperature [14, 92, 94, 95, 97, 100].  14892  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'J Mater Sci  (2018) 53:14885-14906  14893  Columnar ZrO 2  ZrB2-15ZrSi2 (vol%)  ZrO2  ZrB2-15MoSi2 (vol%)  Ta, Zr contained  platelet oxides (TaZr2.75O8)  ZrB2-15TaSi2 (vol%)  Columnar ZrO 2  ZrB2-15WSi 2 (vol%)  SiO2 + ZrO2  Columnar ZrO 2  Dense ZrO 2  Un-oxidized  SiO2 + ZrO2  ZrO2-SiO2 -MoB phase  Un-oxidized  Ta-Zr-O  phase  Coarse ZrO2  Columnar ZrO2  Dense ZrO2  (External surface)  Figure 11 Cross sections and schematics for ZrB2-15 vol% MeSi2 (Me = Zr, Mo, Ta, or W) after oxidation at 1650 °C for 15 min in air  [99].  during the oxidation of ZrB2-TaSi2 at above 1500 °C: Þ þ B2O3 lð Þ þ 1=2O2 g TaSi2 ! TaB2 Þ þ 2SiO2 l; s  sð  sð  ð  Þ  \\x00  \\x01  \\x00  \\x01  ð  Zr; Ta  ÞB2  11=4ZrO2  sð sð  Þ þ O2 g ! TaZr2:75O8 Þ þ 1=2Ta2O5 Þ ! TaZr2:75O8  sð  sð  Þ  temperatures  ð15Þ  ð16Þ  ð17Þ  sð  Þ  The formation of TaB2,  in the form of a ZrB2-TaB2  solid solution ((Zr, Ta)B2),  is effective for improving  the oxidation resistance ever, above 1650 °C,  at  this  temperature. How the oxide layer is composed of  Ta, Zr, and Si containing oxides and a liquid phase is  formed. Furthermore, platelet TaZr2.75O8  formed in  the oxide layer acts as an oxygen-diffusion channel  and Ta2O5, as well as other gaseous species such as  SiO and B2O3,  is evaporated. Thus,  the oxidation of  ZrB2-TaSi2 is accelerated at 1650 °C and the thickness of the oxidized layer after oxidation at 1650 °C reaches * 500 lm.  temperatures  above  \\x0c', 'For  ZrB2-WSi2,  the  oxidation  behavior  is  also  improved by the formation of a SiO2 layer with B2O3 1200-1350 °C in accor and SiO2  after oxidation at  dance with the following reactions: Þ þ 7=2O2 g ! WO3 Þ þ 2SiO2  WSi2  sð  \\x00  \\x01  sð  sð Þ  ð18Þ  Zr; W  ð  ÞB2  sð  Þ þ O2 g  \\x00  \\x01  ! ZrO2  sð  Þ þ WO3  sð  Þ  ð19Þ  However,  the oxidation rate of ZrB2-WSi2 gradually 1500 °C because  increases  at  temperatures  above  columnar ZrO2 is formed in the oxide layer and ZrB2-  ZrSi2  and WO3  are  evaporated as gaseous  species  such as SiO and B2O3. The thickness of the oxide layer formed on ZrB2-WSi2 after oxidation at 1650 °C is * 200 lm and it increases to * 400 lm after oxidation at 1800 °C.  The results of these numerous studies indicate that  the addition of MoSi2 is the most effective method of  improving the oxidation behavior of monolithic ZrB2.  However, SiO2 scales begin to evaporate at temperatures above 1800 °C and an oxide with a high vapor  pressure (MoO3) is formed by the oxidation of MoSi2.  In addition,  the oxidation and related reactions of  ZrB2-MeSi2  composites are quite complex and fur ther  studies on the  reaction mechanisms of MeSi2  containing ZrB2 are required.  Addition of a third phase to improve  the oxidation behavior of binary ZrB2 composites  Various researchers have considered the addition of a  third phase  to  binary ZrB2-based composites  (ZS,  ZrB2-MeSi2,  etc.)  to  improve  their  oxidation resis tance. Opila et al.  [100] added TaSi2 and TaC to ZS  composites and examined their oxidation behavior at and 1927 °C. The liquid phase produced by Ta2O5\\x016ZrO2 was  1627  Ta2O5  and  found  to  cause  a  decrease in the oxidation resistance of the ZS composites at 1927 °C. Hu et al. also added TiB2, TaSi2,  LaB6, CrB2, HfB2, TaB, AlN, and La2O3 to ZS composites and oxidized them at 1800 °C for 1 h [101].  These  experiments  revealed that  the  formation  of  pores, nucleation of cracks, and blistering of the oxide layer occurred after oxidation at 1800 °C for 1 h. They  also noted that only the addition of HfB2  reduced  weight gain during oxidation. Figure 13 shows  the  relationships between the melting points and densi ties of  some UHTCs, oxides, and other  refractories.  Although UHTCs have quite high melting points,  those of  the oxides  evolved by reaction are much  lower. The existence of oxides drives the evolution of  liquid phases during oxidation, as well as material  consumption, because  the  liquid oxide  layers flow  from the surface and the oxygen diffusivity of liquids  is much higher  than that of  solids. Therefore, Zr based or Hf-based composites are considered attrac tive  candidates  for  improving  the  oxidation  Figure 12 Relationship  between  oxidized  layer  thickness  and  oxidation temperature  for ZrB2-15 vol% MeSi2  (Me = Zr, Mo,  Ta, or W) after oxidation for 15 min [99].  Figure 13 Relationship between melting point  and density for  various UHTCs, oxides, and other  refractories [34].  14894  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'resistance of ZS composites perature region ([ 1800 °C).  in the ultra-high-tem Addition of ZrC  Although Zr-based or Hf-based UHTCs are attractive  candidates for improving the oxidation resistance of  ZS  composites,  further  reduction  of material con1800 °C has been  sumption at  temperatures  above  explored. As shown in Fig. 13, although the melting  points of Hf-based UHTCs and HfO2 are higher than  those of Zr-based UHTCs and ZrO2,  the densities of 11 g/cm3  Hf-based UHTCs  are  greater  than  and  much higher than those of Zr-based UHTCs (6-7 g/ cm3). Thus,  light weights and high temperatures are  compatible  for Zr-based UHTCs. Among Zr-based  UHTCs, ZrC has been focused on as an additional  phase for ZS and ZrB2-SiC-ZrC ternary composites  (hereinafter denoted by ZSZ) by various researchers  [57, 58, 68, 102-116], particularly for their potential to  maintain heat resistance and reduce weight  in heat resistant aerospace components. The oxidation reac tion of ZrC can be expressed as follows: ZrC sð Þ þ 3=2O2 g ! ZrO2 Þ þ CO g  \\x00  \\x01  sð  \\x00  \\x01  ð20Þ  ZrC oxidation begins at * 400 °C and the oxidation  rate appears  to be linear with temperature because  porous ZrO2 formed on the ZrC surface does not act  as a barrier  to oxygen diffusion toward the unoxi dized region [117, 118]. Thus,  the oxidation rate of  ZrC is much higher than those of ZrB2 and SiC, which  form B2O3 and SiO2 on their surfaces as barriers to  oxygen diffusion. However, a volume expansion of * 33% accompanies the oxidation of ZrC because the densities of ZrC and ZrO2 are 6.7 and 5.68 g/cm3,  respectively. Thus,  one  of  the motivations  for  the  addition of ZrC to ZS composites  is  to utilize  this  volume expansion to replace material consumed by  the formation of the SiC-depleted layer. Furthermore,  ZrC prevents  the formation of a SiC-depleted layer  owing to the oxidation kinetics of ZSZ composites.  The thermodynamic effect of ZrC on improving ZS  composites is discussed in detail in ‘‘Oxidation below 2000 °C’’ section.  Oxidation behavior of ZSZ composites  Performance evaluation of ZSZ composites in oxida tive environments has been conducted using various  heating systems,  such as  conventional  electric  fur naces,  thermogravimetric  analysis  (TGA)  furnaces,  torch testing,  arc-jet  testing,  and electric  resistance  heating.  In this  section,  the oxidation behavior and  mechanisms of ZSZ composites are reviewed.  Oxidation below 2000 °C  The majority of oxidation studies on ZSZ composites have been carried out at temperatures below 2000 °C  [68, 105-107, 109, 110, 112, 114, 115, 119]. Figure 14a-c  displays a cross section of an oxidized ZSZ composite  obtained by oxidation of ZrB2-16SiC-20ZrC (vol%) heated up to 1500 °C in air, ZrB2-16SiC-20ZrC (vol%)  Figure 14 Cross sections of oxidized ZSZ composites produced temperatures below 2000 °C: a 10 °C/min up to (ZrB2-16SiC-20ZrC (vol%)), 1700 °C for and c 1750 °C for  by oxidation at 1500 °C in  air  b  10 min in air  (ZrB2-16SiC-20ZrC (vol%)),  30 min in air  (ZrB2-20SiC-6ZrC (vol%))  [105, 112, 114].  J Mater Sci  (2018) 53:14885-14906  14895  \\x0c', 'at 1700 °C in air using an oxyhydrogen torch, and 1750 °C in air,  ZrB2-20SiC-6ZrC (vol%)  at  respec tively. As shown in Fig. 14,  the oxidized layers that  form on ZSZ composites during oxidation 2000 °C are mainly SiO2 or SiO2 ? ZrO2  below  layers and  SiC-depleted  layers  (ZrO2 ? ZrB2  layers).  This  is  similar to the results obtained for oxidized ZS com posites. As mentioned in ‘‘Addition of a secondary  phase  to improve  the oxidation behavior of ZrB2’’  section,  the  thickness of  the oxidized layer,  tox,  is  defined as the sum of  the thicknesses of all  the sur face  oxide  scales  (SiO2  layer, ZrO2  layer,  and/or  SiO2 ? ZrO2 layer) and the SiC-depleted layer.  Opeka  et  al.  [119]  oxidized  ZSZ  composites  (ZrB2:ZrC = 10:1, 4:1, or 1:1 and ZrB2:SiC = 2:1) at temperatures up to 1500 °C and a heating rate of 20 °C/min in an Ar/O2  atmosphere using a TGA  furnace. They also oxidized such composites at 1200, 1300, and 1400 °C for 5 h using an electric  furnace  [119]. They concluded that  the weight gain observed  during heating in the temperature range of 700-1500 °C increased with increasing ZrC content.  However,  they  also demonstrated the potential  of  ZSZ composites for use in heat-resistant components  because  the  thickness of  the oxide layer on a ZSZ  composite after isothermal heating at 1200, 1300, or 1400 °C for 5 h is * 100 lm,  regardless of  the oxi dation temperature. In contrast, the thicknesses of the  oxide  layer  on monolithic ZrB2  are  100,  200,  and  400 lm after isothermal heating at 1200, 1300, and 1400 °C, respectively.  Arai et al. [112] oxidized ZSZ composites with four  different  compositions  (SiC = 16  vol% and  ZrB2:  ZrC = 20:64, 34:50, 50:34, or 64:20 (vol%)) in air at temperatures up to 1500 °C using a TGA furnace.  They concluded that an increase in the ZrC content of  a ZSZ composite causes drastic oxidation because to be oxidized at * 450 °C and because  ZrC starts  ZrO2 is formed on ZSZ composites as a result of  the  oxidation of ZrC. ZrO2 does not act as an effective  diffusion barrier for oxygen in the unoxidized region.  The  formation  of  a  porous  SiC-depleted  layer  between  the  oxide  layer  (ZrO2 ? SiO2  layer)  and  unoxidized region was observed after oxidation, as  seen in ZS composites. These ZSZ composites were 1500 °C in  oxidized at  air using  an  infrared (IR)  image furnace combined with an in situ observation  system for  oxidizing  the  surface. As  shown  in  Fig. 15a and b, the results suggest that bubbles form and burst during oxidation at 1500 °C (Fig. 15a) and  that  the diameters  of  the  bubbles  formed on  the  surface  of  the  ZSZ  composites  decrease  with  decreasing ZrC content  (Fig. 15b).  In addition,  the  ZSZ  composite with  the  lowest  content  of  ZrC  (ZrB2:ZrC = 64:20)  exhibited  the  highest  oxidation  resistance because the oxide layer formed on this ZSZ  composite was the thinnest among the ZSZ compos ites oxidized in the study.  A more systematic study of  the effects of compo sition on the oxidation behavior of ZSZ composites  was  reported by Akin and Goller  [68]. They fabri cated ZSZ composites with different  compositions  (ZrB2:ZrC:SiC = 50:40:10,  60:30:10,  60:20:20,  and  50:30:20 (vol%)) and oxidized them at 1400, 1500, and 1600 °C in air and 1500 °C,  for 180 min. At  temperatures of 1400  the weight gain per unit area tended to  increase with increasing temperature and ZrC con tent. ZSZ composites with higher  contents of ZrC  (ZrB2:ZrC:SiC = 50:30:20)  exhibited  lower  weight  gains  than ZSZ composites with lower  contents of  ZrC (ZrB2:ZrC:SiC = 60:20:20) at 1500 and 1600 °C. The results of  temperatures  of  these studies imply  that SiC and ZrC contents of 20-30 vol% moderately improve the oxidation behavior at 1500-1600 °C. The  oxidation  behavior  of  ZSZ  composites  depends  strongly on the proportions of ZrB2, SiC, and ZrC. It  is desirable that the SiC content is lower than the ZrC  content and that  the ZrB2 content  is higher than the  ZrC content.  Wang et al. [105] oxidized ZrB2-20SiC-6ZrC (vol%) composites at 1750 °C for 30 min in air by means of  electric  resistance heating. They indicated that  the  ZSZ composites  formed oxide  scales composed of at 1750 °C. Further SiO2 and ZrO2 on the  surface  more,  controlled oxidation of  the ZSZ composites  proceeded  because  the  relationship  between  the  weight gain of ZSZ composites during oxidation and  time was found to obey a parabolic law.  Kubota et al.  [114] oxidized ZSZ composites with  four  different  compositions  (SiC = 16  vol% and  ZrB2:ZrC = 20:64, 34:50, 50:34, and 64:20 (vol%)) and 1700 °C for  ZS  composites  at  10 min using  an IR  image furnace. The oxidation test was conducted in (2.0 9 103 Pa). They  an air atmosphere at a low PO2  showed that  the thickness of  the oxide layer on the  surface  after  oxidation  increased with  increasing  content of ZrC in the ZSZ composites, regardless of  the PO2 .  In contrast, at a low PO2 ,  the formation of a  SiC-depleted layer was not observed after oxidation  of the ZSZ composites, whereas it was observed after  14896  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'J Mater Sci  (2018) 53:14885-14906  14897  (a)  (b)  ZrB2:SiC:ZrC (vol%)  Asreceived  1200oC  1400oC  1500oC  1500oC2min  After test SEM (BSE)  20: 16: 64  34:16: 50  50: 16: 34  64: 16 :20  500μm  500μm  500μm  500μm  Figure 15  a Formation and bursting of bubbles during oxidation of 64ZrB2-SiC-20ZrC (vol%) at 1500 °C and b in situ observation of  oxidation for ZSZ composites [112].  oxidation of  the ZS composites. Based on the results  an oxyhydrogen torch. Although some spallation of  of  a  thermodynamic  analysis,  they concluded that  the oxide  layer was observed,  the  thickness of  the  preferential oxidation of SiC (Eq. 7)  in comparison  oxide  layer  after  oxidation  of  the  20ZrB2-16SiC-  with ZrB2 (active oxidation of SiC) was accelerated at  64ZrC (vol%) composite was thinner than that of the  low PO2 . However, preferential oxidation of ZrC in  other ZSZ composites. Further, partial  sintering of  comparison with SiC and ZrB2, together with large expansion (* 33%), prevented the  volumetric  for ZrO2  in the  oxide  layer  occurred for  the  20ZrB2-  16SiC-64ZrC (vol%) composite owing to the forma mation of a SiC-depleted layer. They also oxidized the same ZSZ composites at 1700 °C for 10 min using  tion of ZrO2 by oxidation of ZrC being more rapid  than the formation of SiO2 by oxidation of SiC.  In  \\x0c', '14898  J Mater Sci  (2018) 53:14885-14906  confirmed that  the oxidation resistance of ZSZ com posites in oxidation testing with dynamic pressure is  superior  to that of ZS  composites  in terms of  the  durability of the oxide layer under dynamic pressure.  The thickness data for  the oxidized layer on ZSZ  composites are summarized in Fig. 16. For oxidation at temperatures below 2000 °C, an increase in the ZrC  content of a ZSZ composite increases the thickness of  the oxidized layer.  In addition,  the thickness of  the  oxidized layer on a ZSZ composite formed at a low  PO2  is thinner  than that  formed under an air atmo sphere. This oxidation behavior is completely differ ent  from  that  observed  for  ZS  composites.  Figure 16 Thickness of oxidized layer on ZSZ composites as a  function of ZrC content  [68, 105, 112, 114-116].  contrast, the oxide layer formed on the surface of the  ZS composites was removed during a torch test. This  phenomenon was unique  to oxidation testing with  dynamic pressure (a torch test). A series of  studies  Furthermore, the thickness of the oxidized layer on ZSZ composites heated to 800 °C at a heating rate of 10 °C/min is similar to the thickness of the oxidized and 2000 °C. These  obtained at  results  layer  1700  indicate that  the oxidation behavior of ZSZ compos ites depends not only on the ZrC content but also on  the heating rate during oxidation because ZrC starts to be oxidized at * 400 °C. Thus,  the above results  indicate that  the following three conditions are for  (a)  (c)  Temperature increase  (b)  (d)  ZrO2 + SiO2  SiC-dep leted layer  Un-oxid ized reg ion  ZrC content  increase  SiO2  ZrO2 + SiO2  SiC-dep leted layer  Un-oxid ized reg ion  pO2 decrease  ZrO2 + SiO2  SiC-dep leted layer  Un-oxid ized reg ion  ZrO2 + SiO2  Un-oxid ized reg ion  Figure 17  Schematic  illustrations  of  the  oxidation  behavior  of  ZSZ with  16  vol% SiC at  temperatures  below 2000 °C for  ZrB2:ZrC = a 20:64, b 34:50, c 50:34, and d 64:20 (vol%).  \\x0c', 'ZSZ composites 2000 °C:  oxidized  at  temperatures  below  layer are seen upon the addition of ZrC. Thus, ZrC  a ZrC content of  16 vol% or  less,  a ZrB2  addition to form ZSZ composites has proved to be an  content that is higher than the ZrC content, and a SiC  effective  behavior at  strategy for obtaining improved oxidation temperatures above 1600 °C.  Oxidation above 2000 °C  It is quite difficult to conduct above 2000 °C because almost all  oxidation  tests  conventional  fur naces  with  Fe-Cr-Ni  alloy  (  content  that  is lower than the ZrC content.  Figure 17a-d shows  schematic illustrations of  the  oxidation behavior of ZSZ composites at temperatures below 2000 °C. A SiO2 ? ZrO2 layer is formed temperatures above 1600 °C, a SiO2  and at  layer  is  formed on the ZrO2 ? SiO2 layer (Fig. 17a-d). Then,  a SiC-depleted layer  is formed between the ZrO2 ?  SiO2 layer and the unoxidized region. However,  the  SiC-depleted layer disappears as PO2 is decreased (Fig. 17d). At temperatures below 2000 °C (especially 800-1500 °C), a larger amount of ZrC in ZSZ com posites leads to greater weight gain and the forma tion of a thicker oxide layer during oxidation because  the oxidation rate of ZrC is much higher than those of  ZrB2 and SiC (Fig. 17b). Unfortunately,  the effective ness of ZrC for improving the oxidation behavior is confirmed up to 1500 °C. However, (1600-1700 °C),  the disappearance  temperatures  at higher  not  of  the SiC-depleted layer and the formation of a rigid  oxide layer by partial sintering of ZrO2 in the oxide  \\x0c', 'Recently, we constructed a resistant heating system  with a ceramic heater  to oxidize ZS and ZSZ com2000 °C,  posites  at  temperatures  above  which  allowed  us  to  characterize  the  initial  oxidation  behaviors of ZS (20 vol% SiC) and ZSZ composites  (SiC = 16 vol% ZrB2:ZrC = 20:64, 34:50, 50:34, and 64:20 (vol%)) at temperatures above 2000 °C [96, 116].  Figure 18a-d displays cross sections of ZSZ composites oxidized above 2000 °C. As shown in Fig. 18,  a SiC-depleted layer is not formed in ZSZ composites,  but  it  is formed in ZS composites (cf. Figs. 9d, e and  10). ZSZ composites with high contents of ZrC (64  vol%) have only a dense oxide layer, mainly com posed of ZrO2, and their behavior is quite different from that of ZSZ composites formed below 2000 °C  (cf. Fig. 14a-c). Figure 19a-d shows schematic illus trations of  the initial oxidation mechanisms of ZSZ temperatures above 2000 °C.  composites oxidized at  For  all ZSZ composites,  an oxidized layer with a  thickness of  20-30 lm was  formed, but unlike ZS  composites under  the  same  conditions, no SiC-de pleted layer was formed. This result confirmed that  the addition of ZrC to ZS composites prevented the  formation of a SiC-depleted layer during the initial  stage of oxidation, namely short time (10 s) exposure at 2000 °C.  In addition,  the ZSZ composite with the  highest content of ZrC (20ZrB2-16SiC-64ZrC (vol%))  formed a rigid ZrO2 layer on the surface, whereas the  other composites formed a porous ZrO2 layer as the  oxide layer. As Fig. 16 shows,  the thicknesses of  the  oxidized layers on ZS and ZSZ composites are com parable. Thus,  the formation of an oxide layer with out  a  SiC-depleted  layer  and  the  acceleration  of  oxidation in comparison with ZS were only achieved  when ZSZ composites were oxidized at temperatures above 2000 °C. The effect of adding ZrC on improv ing the oxidation behavior of ZSZ composites relative  to that of ZS composites is demonstrated at atures above 2000 °C.  temper The  oxidation behavior  of ZSZ composites was  analyzed based on thermodynamics. Figure 20 dis plays volatility diagrams for ZSZ composites at (a) 1727 °C and (b) 2027 °C. As shown in Fig. 20a and b,  oxidation of ZrC occurs at a much lower PO2  than  Dense  ZrO2 layer  ZrC  SiC  Un-oxidized  region  Pore  ZrB2  v  Dense  ZrO2 layer  Porous  ZrO2 layer  Pore  Un-oxidized  region  (b)  ZSZ50  Dense  ZrO2 layer  Pore  Un-oxidized  region  SiC  ZrB2  (c)  ZSZ34  Partially dense ZrO2 layer Porous  ZrO2 layer  Continuous pore  Un-oxidized  region  (d)  ZSZ20  (a) ZSZ64  Porous  ZrO2 layer  ZrC  Figure 19  Schematic illustrations of the initial oxidation mechanisms of ZSZ composites with ZrC contents of a 64 vol% b 50 vol% c 34 temperatures above 2000 °C [116].  vol% and d 20 vol% oxidized at  14900  J Mater Sci  (2018) 53:14885-14906  \\x0c', 'oxidation of ZrB2 and SiC. At temperatures below 2000 °C, preferential oxidation of ZrC occurs with  simultaneous active oxidation of SiC (existence of SiO  (g) in Fig. 20), as the PO2 required for oxidation of ZrC  and SiC is much lower than that required for oxida tion of ZrB2. The oxidation rate of ZrB2 decreases  with decreasing PO2  (Fig. 5). Thus,  the ZSZ compos ites with the lowest ZrC content have the thinnest  oxide layers. Above 2000 °C,  the PO2  required for oxidation of required below 2000 °C.  ZrC is  greater  than  that  Thus, preferential 2000 °C but  oxidation  of ZrC occurs  below  is  prevented  at  temperatures  above  2000 °C. A layer  containing SiO2  is not  formed, as  shown in Fig. 18, because SiO is removed from the  surface by active oxidation of SiC (and reoxidized out  of  the  system)  [96,  116].  In contrast  to SiC (active  oxidation), ZrC exhibits (* 33 vol%) during  a  large volume  expansion  oxidation, which is  the main  reason for  the  formation of  a dense  layer on ZSZ  composites with higher ZrC contents. Therefore,  the  effect of ZrC on improving ZSZ composites occurs 2000 °C. The  only  after  oxidation above  oxidation  behavior of ZSZ is highly dependent on the heating  rate, ZrC content, oxidation temperature, and hold ing time. To maximize the effect of ZrC addition to  ZS composites for improving the oxidation behavior,  it is desirable to use ZSZ composites at temperatures 2000 °C because ZrC promotes oxidation of below 2000 °C and various  above  ZSZ composites  other  binary ZrB2-based UHTC composites (ZS and ZrB2-  MeSi2) 2000 °C.  also  exhibit  oxidation  resistance  below  Further,  the  optimal  composition of ZSZ  composites varies depending on the intended application at temperatures above 2000 °C. Consequently,  guidelines  for material design are  required for  the  use of ZSZ composites in heat-resistant materials and  thermal protection systems.  Summary and outlook  The characteristics of monolithic ZrB2 were reviewed  in this paper. The problems associated with the use of  monolithic ZrB2 in high-temperature structures were  discussed  based  on  experimental  and  theoretical  analyses  of  its  oxidation behavior. These issues temperatures above 1200 °C. The  become obvious at  addition of SiC is an effective way to improve the  oxidation behavior 1600-1700 °C.  of ZrB2  at  temperatures up to  Efforts  to  improve  the  oxidation  behavior  of  ZrB2-SiC (ZS)  composites were  also  reviewed. Further,  the effectiveness of ZrC addition  to  ZS  composites  to  form ZSZ  composites  on  improving  the  oxidation resistance was described.  Based on this review,  the following conclusions and  recommendations can be made.  The oxidation behavior of ZSZ composites depends  on the test conditions, such as the oxidation temper ature, PO2 , and heating rate. At 1500 °C in  temperatures up to  air,  increasing  the ZrC content  causes  drastic material consumption because ZrC is significantly oxidized. At temperatures below 2000 °C, ZSZ  Figure 20 Volatility diagrams for ZSZ composites at a 1727 °C and b 2027 °C [114, 116].  J Mater Sci  (2018) 53:14885-14906  14901  \\x0c', 'composites  (ZrB2:SiC:ZrC = 64:16:20  (vol%))  show  less material consumption, as do ZS composites oxidized at 1700 °C. Moreover, at a low (2.0 9 103 Pa), a SiC-depleted layer is not formed on  PO2  ZSZ composites, whereas it is formed on ZS composites during oxidation at 1700 °C. These observa tions  imply  that  the  oxidation  behavior  of  ZS  composites at low PO2 is probably improved by the addition of ZrC, even at temperatures below 2000 °C.  Consequently,  the addition of ZrC to ZS composites  is quite effective in the case of oxidation at temperatures above 2000 °C because preferential oxidation  of ZrC over SiC prevents the formation of a SiC-de pleted layer. In addition, even at temperatures below 2000 °C, ZrC has been confirmed to effectively improve the oxidation behavior of ZSZ at 1700 °C in  low PO2 atmospheres because less PO2  is required to  oxidize ZrC than SiC and the  formation of  a SiC depleted layer is also prevented by the oxidation of  ZrC.  Oxidation studies of ZSZ composites conducted at 2000 °C have  temperatures  above  shown  that  the  thickness of the oxidized layer on ZSZ composites is  comparable to that on ZS composites. A SiC-depleted  layer  is not  formed on ZSZ composites during oxi dation testing, but  it  is  formed on ZS  composites  under the same conditions.  In addition, a ZSZ com posite  with  64  vol%  ZrC  [20ZrB2-16SiC-64ZrC  (vol%)]  formed a dense ZrO2  layer as  the oxidized  layer, whereas ZSZ composites with 20, 34, and 50  vol% ZrC formed porous ZrO2 layers. Thus, the effect  of ZrC on improving the oxidation behavior  (espe cially the formation of a dense layer) was confirmed temperatures above 2000 °C. Further  by oxidation at  studies are required to ensure the good performance 2000 °C  of ZSZ composites  at  temperatures  above  because the holding time at  the maximum tempera ture was limited in the studies conducted to date. The  difficulty  of  testing  at  such  high  temperatures  remains an obstacle to fully understanding the oxi dation behaviors of  such materials. Further  investi gation using arc-wind tunnel testing over a long time  period is in progress.  Compliance with ethical standards  Conflict of  interest  The authors declare that  they  have no conflict of  interest.  References  [1]  Fahrenholtz WG, Hilmas GE, Chamberlain AL, Zimmer mann JW (2004) Processing and characterization of ZrB2 based  ultra-high  temperature  monolithic  and  fibrous  monolithic  ceramics.  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},{
  "_id": 178,
  "PDF": "Oxidation of ZrB2 powder in the temperature range of 650–800°C.pdf",
  "Text": "['Journal of Alloys and Compounds 471 (2009) 502-506  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j a l l c o m  Oxidation of ZrB2 powder in the temperature range of 650-800  C  Wei-Ming Guo, Guo-Jun Zhang ∗ , Yan-Mei Kan, Pei-Ling Wang  State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Shanghai 200050, China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 23 February 2008 Received in revised form 1 April 2008 Accepted 3 April 2008 Available online 19 May 2008  Keywords:  Kinetics Microstructure Oxidation  1.  Introduction  The isothermal oxidation of ZrB2 powder was carried out in the range of 650-800   C in a flowing air using thermogravimetric analysis (TGA). The evolution of the phase characterization and morphology of ZrB2 powder oxidized at 700   C for varying durations was studied by X-ray diffraction (XRD) and scanning electron microscopy (SEM), respectively. The excellent fit of TG curves by multiple-law model suggests that the oxidation of ZrB2 powder in air follows para-linear kinetics, and based on the fitted results, oxidation mechanisms can also be obtained. The reaction product of ZrB2 powder with oxygen is metastable tetragonal ZrO2 at 700   C, and the tetragonal phase transforms to the monoclinic phase with oxidation. The oxidation of ZrB2 powder is associated with surface microcrack formation, which is attributed to volume expansion resulting from oxidation of ZrB2 to tetragonal ZrO2 and tetragonal-to-monoclinic phase transformation of ZrO2 . In the last stage of oxidation, each ZrB2 particle breaks into fragments. © 2008 Elsevier B.V. All rights reserved.  Many transition-metal diborides have attracted considerable attention owing to the excellent combination of physicochemical properties such as high melting temperature, high hardness, and high thermal conductivity. ZrB2 is one of these diborides and a candidate for use in the thermal protection systems and scramjet engine components for hypersonic flight vehicles, as well as high temperature electrodes, molten metal containment systems, and incinerators, because of its good resistance to ablation at high temperatures and thermal-shock resistance in addition to the attractive aforementioned properties [1,2]. When ZrB2 is exposed to oxidizing environments at high temperatures, it oxidizes, which will degrade its properties. Thus, a large number of studies on oxidation of ZrB2 and its composites in air have been carried out in order to better understand the oxidation mechanisms and enhance oxidation resistance [3-12]. The reaction of ZrB2 with O2 is as follows: ZrB2 + 5 2 O2 (g) → ZrO2 + B2O3 (l) (1) B2O3 is known to have a low melting point of 450  C and high vapor pressure, so it readily vaporizes at high temperatures. B2O3 (l) → B2O3 (g)  (2)  Researchers have studied the products resulting from the oxidation of ZrB2 , and discovered that the oxide scale is composed of  ∗  Corresponding author. Tel.: +86 21 52411080; fax: +86 21 52413122. E-mail address: gjzhang@mail.sic.ac.cn (G.-J. Zhang).  0925-8388/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.04.006  an external glassy B2O3 layer and an underlying porous ZrO2 layer filled with B2O3 . The glassy B2O3 film acts as an effective barrier against oxygen diffusion, and the porous ZrO2 layer does not offer temperatures below 1100  C, a parabolic effective protection. At behavior was reported for the oxidation of ZrB2 due to the presence of a continuous layer of B2O3 [3,8-10]. B2O3 rapidly vaporizes at temperatures above 1100  C, thus reducing the effectiveness of the diffusion barrier. Para-linear kinetics have been observed in the temperature range of 1100-1400  C [4,8-10]. Because the removal of the B2O3 layer above 1400  C leaves behind a porous ZrO2 scale, ZrB2 exhibits rapid linear oxidation kinetics at these temperatures [9,10]. It has been almost 50 years since researchers started to explore the oxidation of ZrB2 . However, previous investigations have mainly focused on the oxidation of bulk materials, oxidation mechanisms of ZrB2 powder as well as phase composition and microstructures that evolve during oxidation have not been studied in detail. In this paper, the oxidation behavior of ZrB2 powder from 650 to 800  C in air was investigated. Firstly, the isothermal oxidation kinetics and mechanisms were studied by multiple-law model. Then the evolution of phase characterization and morphology of the samples before and after oxidation was described.  2.  Experimental  The ZrB2 powder employed in this study was a commercial product (Kojundo Chemical Laboratory Co. Ltd, Saitama, Japan). The surface area of the sample was 0.3042 m2 /g, as measured by the BET method. Oxidation tests were carried out by a Netzsch STA449C simultaneous thermal analyzer under isothermal conditions. The samples were heated to the desired temperature (650, 700, 750, and 800   C) for the isothermal oxidation in flowing nitrogen (99.99%, 50 ml/min). The fast heating stage with a heating rate of 50   C/min prior to the isothermal period was applied  \\x0c', 'W.-M. Guo et al. / Journal of Alloys and Compounds 471 (2009) 502-506  503  Fig. 1. Thermogravimetric curves of ZrB2 powder oxidized in air at different temperatures: (a) 650   C; (b) 700   C; (c) 750   C; and (d) 800   C.  Fig. 2. Square of mass gain versus time for ZrB2 powder oxidized in air at different temperatures: (a) 650   C; (b) 700   C; (c) 750   C; and (d) 800   C.  to minimize oxidation effects before reaching the target temperatures. When the desired temperature was reached, the sample was held at that temperature for about 5 min before the gas was switched from nitrogen to air. Roughly 60 mg of ZrB2 powder was placed into an alumina crucible and oxidized in flowing air (50 ml/min). Due to the low oxidation temperatures, no obvious interaction between the sample and the crucible was observed. Phase composition and morphology of the samples before and after oxidation were studied by X-ray diffraction (XRD) with Cu K␣ radiation and scanning electron microscopy (SEM), respectively.  3. Results and discussion  3.1. Oxidation kinetics and mechanisms  TG curves of ZrB2 powder oxidized in air at 650-800  C are shown in Fig. 1, where w is the mass gain per unit surface area and  Fig. 3. Multiple-law modeling of TG curves of ZrB2 powder oxidized in air at different temperatures: (a) 650   C; (b) 700   C; (c) 750   C; and (d) 800   C.  \\x0c', '504  W.-M. Guo et al. / Journal of Alloys and Compounds 471 (2009) 502-506  t is oxidation time. With increasing temperature, it is obvious that the weight gain increases at identical time. It can be seen that the TG curve approaches linear behavior for 650  C, whereas TG curves clearly show deviation from linear behavior with the increase of temperature. Fig. 2 presents the square of mass gain per unit surface area as a function of oxidation time for ZrB2 powder. It is found that the square of mass gain per unit surface area is a nonlinear function of oxidation time, which indicates that oxidation of ZrB2 powder deviates from parabolic kinetics. Therefore, a simple parabolic law is unable to describe the kinetic behavior. In order to give a complete analytical description, the TG curves were fitted according to a multiple-law model including a linear and parabolic term, developed by Nickel [13]. The mass change per unit area w is expressed as a function of time, according to the following equation:  √  t  w = A + Klin t + Kpar  (3)  and Kpar  where A is a constant, Klin is linear rate constant is parabolic rate constant. The fit allows for the quantification of the individual contributions of the multiple-law model and accounts for para-linear kinetics. Using multiple linear-regression analysis, multiple-law modeling of TG curves of ZrB2 powder oxidized in air at different temperatures was performed, with the results illustrated in Fig. 3. The K parameters and R2 values corresponding to each fitting curve are reported in Table 1. The excellent fit of the solutions confirms that oxidation of ZrB2 powder follows paralinear kinetics in such a temperature range (R2 > 0.99). As seen in Fig. 3, the dominating term is the parabolic one in the range of 650-800  C, accounting for oxygen diffusion in oxide scale. From 650 to 700  C, the linear term is positive, which is related to mass gain due to chemical reaction between ZrB2 and O2 at the In the range of 750-800  C, diboride-oxide interfaces. the linear term is negative, which is associated with mass loss due to vaporization of B2O3 . It can therefore be inferred that oxidation of ZrB2 powder in air is governed by oxygen diffusion and chemical reaction at 650-700  C, but controlled by oxygen diffusion and the evaporation of B2O3 at 750-800  C. Even though the vapor pressure of B2O3 may be low in the temperature range of 750-800  C, oxidation kinetics revealed that the evaporation of B2O3 should not be neglected during the oxidation of ZrB2 powder. This deduction can be also supported by the study on oxidation of AlN-SiC-ZrB2 composites reported by Brach et al. [14]. In the range 700-900  C, Brach et al. [14] indicated the oxidation of AlN-SiC-ZrB2 composites followed mixed para-linear kinetics, where the linear term was negative, accounting for mass loss due to vaporization of B2O3 .  3.2. Phase characterization and morphology of the oxidized samples  Fig. 4 shows the XRD patterns of ZrB2 powder oxidized at 700  C for varying durations. Before the exposure, only ZrB2 peaks are detected as shown in Fig. 4a. In the initial stage of oxidation, tetragonal ZrO2 peaks appear (Fig. 4b). With increasing oxidation time,  Table 1 Kinetic parameters of the oxidation fitting curve, relatively to parabolic and linear contributions, and R2 (fit goodness) values.  Oxidation temperature (  C)  Kpar  (mg cm−2 min−0.5 )  Klin  (mg cm−2 min−1 )  R2  650 700 750 800  0.00192 0.00536 0.02068 0.03405  0.00013 −0.00040 0.00007 −0.00122  0.99852 0.99947 0.99744 0.99886  Fig. 4. X-ray diffraction patterns of ZrB2 powder oxidized at 700   C for varying durations: (a) no oxidation; (b) 0.5 h; (c) 4 h; and (d) 15 h.  monoclinic ZrO2 peaks also appear, while intensities of ZrB2 peaks simultaneously decrease, as shown in Fig. 4c. When oxidation time is increased to 15 h at 700  C, the XRD patterns show monoclinic ZrO2 as a major phase, a minor phase of tetragonal ZrO2 , and the disappearance of the ZrB2 phase (Fig. 4d). Generally, the stable polymorph of zirconia at room temperature and atmospheric pressure is monoclinic, which transforms at 1170  C to tetragonal and then at 2370  C to cubic structure. Previous researchers [15] have investigated the effect of the crystallite size on the phase transformation and pointed out that the phase transformation was closely related to the growth of the crystallite size. When the particle size reaches nano-scale, a metastable tetragonal ZrO2 can also exist at room temperature, which is so-called nanosize effect of particle phase [16]. So, the reaction of ZrB2 powder with oxygen can generate metastable tetragonal ZrO2 instead of stable monoclinic phase at 700  C. Broad ZrO2 peaks also validate this deduction in Fig. 4b. However, the tetragonal crystallites grow during annealing. Upon reaching a critical diameter, the tetragonal phase transforms immediately to the monoclinic phase. Therefore, monoclinic ZrO2 peaks appear with oxidation. The increase of holding time will induce more phase transform of ZrO2 , and as a result, the oxidized samples after 15 h contains major monoclinic ZrO2 phase with minor tetragonal ZrO2 phase. Morphology changes occurring to the ZrB2 powder as a result of oxidation were examined using SEM. Fig. 5 is a montage of SEM images illustrating the progression of oxidation at 700  C. The asreceived ZrB2 powder consists of irregularly sized particles with mainly columnar shape, as shown in Fig. 5a. The first visible evidence of appreciable oxidation of the samples is the appearance of circumferential microcracks at the edge of the ends of a number of cylinders, for example, on the arrow sites in Fig. 5b. At the beginning of oxidation, the particles keep their original shape, while the surfaces become rough. With further oxidation, microcracks at the surface of the oxide layers running parallel to the central axes of the cylinders are also observed (indicated by arrows in Fig. 5c). At the same time, the oxide scale at the ends of certain columnar particles flake off. After 15 h, every ZrB2 particle has broken into several fragments (Fig. 5d). It is assumed that oxidation products have theoretical density. Therefore, the oxidation from ZrB2 to tetragonal ZrO2 is accompanied by 9% volume expansion, based on the densities of 6.09 and  \\x0c', 'W.-M. Guo et al. / Journal of Alloys and Compounds 471 (2009) 502-506  505  Fig. 5. SEM photographs of ZrB2 powder oxidized at 700   C for varying durations: (a) No oxidation; (b) 0.5 h, arrows indicated the microcracks of the ends of columnar particles; (c) 4 h, arrows indicated microcracks of the oxide layers running parallel to the central axes of columnar particles; and (d) 15 h.  6.10 g/cm3 for ZrB2 and tetragonal ZrO2 , respectively. In addition, the tetragonal-to-monoclinic phase transformation is accompanied by 5% volume expansion [17]. It is thus understandable that oxidation of ZrB2 powder is associated with the generation of surface microcracks. Compared with a flat surface, greater stresses induced by volume expansion exist in the edges of columnar ZrB2 particles. Therefore, microcracks firstly emerge at these edges. With the progress of oxidation, the volume expansion which develops in the oxide layer increases and eventually breaks the particles into fragments.  4. Conclusions  The oxidation of ZrB2 powder follows para-linear kinetics in air at 650-800  C, where the dominating term is the parabolic one, accounting for oxygen diffusion in the oxide scale. From 650 to 700  C, the linear term is positive, which is related to mass gain due to chemical reaction between ZrB2 and O2 at the diboride-oxide interfaces. From 750 to 800  C, the linear term is negative, accounting for mass loss due to vaporization of B2O3 . Therefore, even though the vapor pressure of B2O3 is low in the in the range of 750-800  C, the effect of evaporation of B2O3 should not be neglected during the oxidation of ZrB2 powder. The reaction product of ZrB2 powder with oxygen is metastable tetragonal ZrO2  at 700  C, and the tetragonal phase transforms to the monoclinic phase with oxidation. The oxidation of ZrB2 powder is associated with surface microcrack formation, which can be attributed to volume expansion resulting from oxidation of ZrB2 to tetragonal ZrO2 and tetragonal-to-monoclinic phase transformation of ZrO2 . Finally, in the last stage of oxidation, each particle breaks into fragments.  Acknowledgement  Financial support from the Chinese Academy of Sciences under the Program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program), the National Natural Science Foundation of China (No. 50632070 and No. 50602048) is gratefully appreciated.  References  [1] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Ceram. Soc. 89 (2006) 450-456. [2] F. Monteverde, A. Bellosi, S. Guicciardi, J. Eur. Ceram. Soc. 22 (2002) 279-288. [3] A.K. Kuriakose, J.L. Margrave, J. Electrochem. Soc. 111 (1964) 827-831. [4] W.C. Tripp, H.C. Graham, J. Electrochem. Soc. 118 (1971) 1195-1199. [5] J.B. Berkowitz-Mattuck, J. Electrochem. Soc. 113 (1966) 908-914. [6] R.J. Irving, I.G. Worsley, J. Less-Common Met. 16 (1968) 103-112. [7] F. Monteverde, A. Bellosi, J. Electrochem. Soc. 150 (2003) B552-B559. [8] D. Sciti, M. Brach, A. Bellosi, Scripta Mater. 53 (2005) 1297-1302.  \\x0c', '506  W.-M. Guo et al. / Journal of Alloys and Compounds 471 (2009) 502-506  Fahrenholtz, G.E. Hilmas,  [9] A. Rezaie, W.G. 3240-3245. [10] W.G. Fahrenholtz, J. Am. Ceram. Soc. 90 (2007) 143-148. [11] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, J. Eur. Ceram. 2495-2501. [12] T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerans, Acta Mater. 55 (2007) 5999-6010.  Ceram.  (2007)  Soc.  89  (2006)  J. Am.  Soc.  27  [13] K.G. Nickel, in: K.G. Nickel (Ed.), Corrosion of Advanced Ceramics, Kluwer Academic Publishers, Dordrecht, NL, 1994, pp. 59-72. [14] M. Brach, D. Sciti, A. Balbo, A. bellosi, J. Eur. Ceram. Soc. 25 (2005) 1771-1780. [15] N. Igawa, T. Nagaski, Y. Ishii, K. Noda, H. Ohno, Y. Morii, J.A. Fernandez-Baca, J. Mater. Sci. 33 (1998) 4747-4758. [16] X.M. Liu, Z.F. Yan, Chin. J. Struct. Chem. 25 (2006) 424-432. [17] S. Shukla, S. Seal, Int. Mater. Rev. 50 (2005) 45-64.  \\x0c']"
},{
  "_id": 179,
  "PDF": "Oxidation of ZrB2- and HfB2-based ultra-high temperature ceramics Effect of Ta additions.pdf",
  "Text": "['J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5969 - 5977  Oxidation of ZrB2and HfB2-based ultra-high temperature ceramics: Effect of Ta additions  ULTRA-HIGH TEMPERATURE CERAMICS  E . O P I L A , S . L E V I N E NASA Glenn Research Center, Brookpark, OH 44135, USA E-mail: Elizabeth.J.Opila@grc.nasa.gov  J . L O R I N C Z Ohio University, Athens, OH 45701, USA  Several compositions of ZrB2 and HfB2 -based Ultra-High Temperature Ceramics (UHTC) were oxidized in stagnant air at 1627 C in ten minute cycles for times up to 100 min. These compositions include: ZrB2 -20 vol% SiC, HfB2 -20 vol% SiC, ZrB2 -20 vol% SiC-20 vol% TaSi2 , ZrB2 -33 vol% SiC, HfB2 -20 vol% SiC-20 vol% TaSi2 , and ZrB2 -20 vol% SiC-20 vol% TaC. The weight change due to oxidation was recorded. The ZrB2 -20 vol% SiC-20 vol% TaSi2 composition was also oxidized in stagnant air at 1927 C and in an arc jet atmosphere. Samples were analyzed after oxidation by X-ray diffraction, field emission scanning electron microscopy, and energy dispersive spectroscopy to determine the reaction products and to observe the microstructure. The ZrB2 -20 vol% SiC-20 vol% TaSi2 showed the lowest oxidation rate at 1627 C, but performed poorly under the more extreme tests due to liquid phase formation. Effects of Ta-additions on the oxidation of the C(cid:2) 2004 Kluwer Academic Publishers diboride-based UHTC are discussed.  1.  Introduction  Preliminary results for the oxidation of ZrB2 -20 vol% SiC-20 vol% TaSi2 [1] showed improved oxidation resistance relative to the baseline material ZrB2 -20 vol% SiC. This improved behavior was attributed to the addition of Ta to the system. One possibility for the improved oxidation resistance, explored by Talmy et al. [2], is that Ta2O5 in the borosilicate glass causes liquid immiscibility and phase-separated glasses of higher viscosity and lower permeability to oxygen. Talmy et al. found TaB2 additions to ZrB2 -20 v/o SiC were more effective in improving oxidation resistance at temperatures between 1200 and 1400 C than additions of other group IV-VI transition metal borides, including Cr, Nb, Ti, and V. No significant improvement in oxidation resistance at temperatures of 1500 C was found for any additions. Another possible explanation for this improved behavior is that Ta additions result in substitution of Ta on the Zr site in ZrO2 , reducing the concentration of oxygen vacancies in the ZrO2 per the following doping reaction given in standard Kroger-Vink notation:        Ta2O5 + V  ••  O  2 ZrO2−→ 2 Ta  •  Zr  + 5 OO  (1)  The resultant lower concentration of oxygen vacancies decreases oxygen transport through the growing oxide scale, and thus lowers the oxidation rate of tantalumcontaining UHTC materials. There is some precedence for Ta2O5 acting as a vacancy suppressor in ZrO2 -based materials [3, 4]. Ionic conductivity in 90 m/o ZrO2 -10 m/o Y2O3 decreased with Ta2O5 additions up to 10  0022-2461  C(cid:2) 2004 Kluwer Academic Publishers  mol% [4]. Greater Ta2O5 additions did not result in any further decrease in ionic conductivity. The purpose of this work is to explore effects of Taadditions to UHTC in hopes of improving oxidation properties for this class of materials to enable their use in space transportation leading edge applications for short times at very high temperatures. This paper represents a preliminary investigation on the oxidation resistance of Ta-containing UHTC.  2. Experimental procedure  ZrB2 and HfB2 -based UHTC were prepared from powders described in Table I. Six compositions, their designation, and processing history are summarized in Table II. The powders were mixed, and then milled using Si3N4 media in hexane in a Si3N4 mill for 24 h. The powders were vacuum hot pressed using a graphite die. Further details can be found in reference [1]. Sample coupons of 2.5 cm × 1.2 cm × 0.3 cm were machined from hot pressed plates. The coupons were ultrasonically cleaned in detergent (Micro-90, International Products Corporation), de-ionized water, acetone, and alcohol prior to oxidation. X-ray diffraction and Energy Dispersive Spectroscopy of as-machined and cleaned coupons indicated the desired phase assemblage was formed in all cases except the ZSTS and ZSTC material. Analysis indicated ZSTS contained ZrB2 , SiC, TaSi2 and possibly a minor amount of a (Zr, Ta) boride solution. ZSTC contained ZrB2 , SiC, and (Zr, Ta) boride and carbide solution phases. Initial surface areas and sample weights were recorded.  5969  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E I  Powders used for UHTC sample preparation  Material  Source  Particle size  Purity (%)  ZrB2 HfB2 SiC  Cerac  <10 µm −325 mesh <5 µm 80% −325 mesh −325 mesh  99.5  Cerac  99.5  H.C. Starck  >99.9  TaSi2 TaC  Cerac  >99.9  Cerac  99.5  T A B L E  I I  Summary  of UHTC compositions,  designations,  and  processing  Designation:  oxidation  Density Density (g/cm3 ) (%)a  Composition  temp.  Processing  ZrB2 -20 vol% SiC  ZS  2000     C, 10 ksi,  5.27  95.5  2 h  HfB2 -20 vol% SiC  HS  2000     C, 10 ksi,  9.60  97.0  2 h  ZrB2 -20 vol% SiC-20 vol% TaSi2 ZrB2 -20 vol% SiC-20 vol% TaSi2 ZrB2 -33 vol% SiC  ZSTS:  1600     C, 10 ksi,  5.92  97.7  1627     C  2 h  ZSTS:  1600     C, 10 ksi,  6.04  99.7  1927 ZS+     C  2 h  2000     C, 10 ksi,  5.04  97.5  2 h  ZrB2 -20 vol% SiC-20 vol% TaC  ZSTC  2000     C, 10 ksi,  7.12  98.9  2 h  HfB2 -20 vol% SiC-20 vol% TaSi2  HSTS  1700     C, 10 ksi,  9.13  100.0  2 h  aBased on rule-of-mixtures.  T A B L E  I I I XRD results and parabolic oxidation rates  for UHTC  materials oxidized in air at 1627     C. XRD results in bold indicate major  phases present  kp  Composition  Designation XRD results  (mg2 /cm4 h)  ZrB2 -20 vol% SiC HfB2 -20 vol% SiC ZrB2 -20 vol% SiC-20 vol% TaSi2 ZrB2 -33 vol% SiC ZrB2 -20 vol% SiC-20 vol% TaC  ZS  ZrO2 (m), ZrO2 (c)  10.94  HS  HfO2 (m)  2.52  ZSTS  ZrO2 (m), ZrO2 (c)  0.29  ZS+ ZSTC  ZrO2 (m), ZrO2 (c) ZrO2 (t), ZrO2 (m)  NA  NA  HfB2 -20 vol% SiC-20 vol% TaSi2  HSTS  HfO2 (m), HfO2 (c), HfSiO4  5.73  (m) = monoclinic, (c) = cubic, (t) = tetragonal.  Samples were oxidized at 1627 C in stagnant air in bottom loading box furnaces, either a DelTech, Inc. zirconia element furnace or a CM Inc. MoSi2 element furnace. Oxidation of ZrB2 -20 vol% SiC-20 vol% TaSi2 was also conducted in the DelTech zirconia element furnace at 1927 C. Three coupons were loaded into a slotted calcia-stabilized zirconia firebrick setter (98.7% purity). Two lines of contact existed between the setter and sample. The coupons were oxidized for ten minute cycles. A coupon was removed after one cycle, five cycles, and ten cycles. A maximum oxidation time of 100 min was achieved. Post-oxidation weight changes were recorded where possible. Some coupons stuck to the zirconia setter due to extensive glass formation. Optical macrographs were taken of all coupons after oxidation. The sample surfaces were analyzed by X-ray Diffraction (XRD), Field Emission Scanning Electron Microscopy (FESEM), and Energy Dispersive Spec       troscopy (EDS) to determine phases present before and after oxidation as well as oxide microstructures. Several materials were also examined by FESEM and EDS in cross-sections that were prepared by non-aqueous cutting and polishing procedures. The ZrB2 -20 vol% SiC-20 vol% TaSi2 composition was also tested in an arc jet in the Interactive Heating Facility at NASA Ames Research Center. Details can be found in Reference [5]. The sample reached temperatures between 1800 and 1960 C in a 10 minute exposure at a stagnation pressure of 0.07 atm and a nominal heat flux of 350 W/cm2 .     3. Results  3.1. Oxidation at 1627 C in air  Macrographs of the coupons after oxidation are shown in Fig. 1a through f. X-ray diffraction results are summarized in Table III. In all cases, ZrO2 or HfO2 was the major phase detected on ZrB2 and HfB2 -based materials respectively. HfSiO4 was only detected on HSTS. Plots of the specific weight change vs. oxidation time for coupons exposed at 1627 C in stagnant air are shown in Fig. 2. All compositions except ZSTC showed evidence of the formation of a protective scale, that is, the oxidation rate slowed with time. The results for the remaining compositions are plotted as specific weight change versus square root of time in Fig. 3. The fit of the data to a straight line on this type of plot indicates parabolic oxidation, that is, the oxidation rate is limited by transport of oxidant across the growing oxide scale. The slope of such a line is equivalent to the square root of the parabolic rate constant, kp . From Fig. 3, it can be seen that the results for ZS, HS, ZSTS, and HSTS show reasonably good fits to parabolic kinetics. Parabolic rate constants are reported in Table III for these four compositions. While these numbers can be used to make semi-quantitative comparisons of oxidation rate, they are not strictly comparable because different compositions of reaction products form on each sample type. A more quantitative comparison will be made in future reports by comparing rate constants derived from recession vs. time measured from cross-sections of oxidized samples. FESEM results are shown in Figs 4 through 7. Figs 4 and 5 show the cross-section of oxidized ZS baseline material and ZSTS respectively. Note the much thinner oxide scale on the ZSTS material which is consistent with the low oxidation rate as measured by weight change. Porosity, due to SiC depletion by active oxidation, is observed near the oxide/matrix interface for the ZS material shown in Fig. 4. Fig. 6 shows a surface micrograph of ZSTS. Note the evidence of glass immiscibility. EDS indicates the phase of medium contrast contains Al and Mg impurities. It is not clear that any Ta is present in the glass phases. EDS indicates that Ta may be present in the ZrO2 phase, however, this must be confirmed by another analytical technique. Fig. 7 shows a surface micrograph of the ZSTC material. The porosity in the oxide scale is clearly visible. The non-protective nature of this oxide scale is     5970  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  (a)  (c)  (e)  (b)  (d)  (f)  Figure 1 UHTC oxidized at 1627     C in air for 1, 5, and 10 ten-minute cycles: (a) ZS, (b) ZSTS, (c) ZS+, (d) ZSTC, (e) HS, and (f) HSTS.  consistent with the observed rapid linear weight change kinetics.  3.2. Oxidation at 1927 C     Macrographs of the ZSTS coupons after oxidation at 1927 C in stagnant air are shown in Fig. 8. After five oxidation cycles, the samples were bonded to the setter by liquid oxide phase formation. Weight change measurements were not possible. Fig. 8c was taken after the sample was broken during an attempt to remove the coupon from the setter. The cross-section shows a gap between the scale and the substrate, showing that rapid substrate consumption occurred. XRD of the oxide scale showed monoclinic ZrO2 as the major phase  and cubic ZrO2 as a minor phase. XRD of the substrate with the scale removed gave the same results. In contrast, evidence of Ta2O5 as well as possible zirconium tantalate formation on the substrate below the ZrO2 scale was found by EDS analysis as shown in Fig. 9. However, this identification is complicated by the fact that the Ta Mα peak (1.709 keV) has almost complete overlap with the Si Kα peak (1.739 keV). Ta identification was made primarily by observing the low intensity Ta Lα peak at 8.145 keV. In addition, phases with high Si content would have much darker contrast than those with high Ta content. Wavelength dispersive spectroscopy (WDS) is planned to address the peak overlap problem as well as to resolve the discrepancy between XRD and EDS results. Fig. 10 shows FESEM images of  5971  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 2 Specific weight change for all UHTC compositions oxidized at 1627     C in air.  Figure 3 Fit of specific weight change data to parabolic kinetics for UHTC compositions with SiC additions indicating protective oxide formation at C in air, however, ZS+ does not show good agreement with parabolic kinetics.  1627     the oxide scale outer surface as well as the underside of the intact oxide scale that was removed. The dark silicacontaining phase is observed on the outer surface but only in much smaller amounts on the underside of the oxide. The bright phases on both the outer and underside of the scale were identified as ZrO2 and zirconium tantalate.  4. Discussion     A comparison of the oxidation rate of the ZS and ZSTS compositions at 1627 C in air clearly showed the benefit of TaSi2 additions toward improving the oxidation resistance of ZrB2 -based UHTC materials [1]. These initial results led to a series of questions that we have answered to varying degrees by the work described here.  3.3. Arc jet testing  Fig. 11 shows a macrograph of the 2.54 cm diameter button of ZSTS after arc jet exposure for 10 min. The scale composition and morphology are similar to that observed after oxidation at 1927 C (Fig. 10, left). Again, evidence of liquid phase formation is observed. Similarly, Ta2O5 was tentatively identified below the oxide scale by EDS [4].     5972  4.1.  Is Ta or Si in TaSi2 responsible for the improved oxidation resistance of ZSTS?  The ZS+ composition was fabricated to answer this question. The ZS+ composition contains the same amount of ZrB2 as ZSTS and the same amount of Si (in the form of SiC) as ZSTS. If the increased oxidation resistance of ZSTS relative to ZS was due to the increased amount of Si, and therefore more SiO2  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 4 ZS cross-section after oxidation in air at 1627     C for 100 min. Arrow shows oxide thickness. Representative areas of each phase are labeled.  Figure 5 ZSTS cross-section after oxidation in air at 1627     C for 100 min. Arrows show oxide thickness. Representative areas of each oxide phase  are labeled. The unoxidized ZSTS below the scale contains three contrasting phases: dark SiC; medium ZrB2 and (Zr, Ta)B2 , bright TaSi2 .  is expected that ZS+ formation on oxidation, then it would have oxidation resistance superior to that of ZS. In fact, the opposite occurred. Fig. 1c shows the increased glass formation of this material, while Fig. 2 shows the ZS+ has a higher oxidation rate than both ZS and ZSTS. This result supports the conclusion that Ta additions are responsible for improving the oxidation resistance of ZrB2 -based materials.  4.2. Does the benefit of TaSi2 additions observed at 1627 C extend to higher temperatures?  Clearly the TaSi2 additions to ZS did not provide superior oxidation resistance at temperatures near 1900 C,           but in fact the opposite. Liquid phases were formed in both the 1927 C furnace oxidation and the arc jet tests. This can be expected based on the phase diagram for the ZrO2 -Ta2O5 system [6]. The melting point of Ta2O5 is 1785 C for the low-temperature phase and 1872 C for the high-temperature phase [7]. The observed liquid could be due to melting of Ta2O5 . However, a zirconium tantalate phase, called phase V, may also be responsible for this liquid phase formation. Phase V is of variable composition and has been described as Ta2O5 ·6ZrO2 [8], Ta2O5 ·5.5 ZrO2 [9], or containing 11 to 17% Ta2O5 in ZrO2 [10]. The melting point of phase V is uncertain and has only been defined to be higher than 1870 C [8]. It is difficult to unambiguously distinguish phase V from cubic zirconia by XRD. The introduction of        5973  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 6 Surface morphology of ZSTS after oxidation in air at 1627  medium phase, silicate glass with impurities; dark phase, SiO2 .     C for 100 min showing evidence of glass immiscibility. Bright phase, ZrO2 ;  Figure 7 Surface morphology of ZSTC after oxidation in air at 1627     C for 100 min.  Figure 8 ZSTS after oxidation at 1927     C in stagnant air: (a) as-received and 1 cycle, (b) 5 ten-min cycles in setter, and (c) 5 ten-min cycles.  5974  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 9 ZSTS substrate (loose scale has been removed) after exposure in stagnant air at 1927     C for 5 ten-min cycles. EDS results on flatter regions  of the bright phase are consistent with Ta2O5 , while EDS on the bright thin aligned grains is consistent with zirconium tantalate. The dark phase is identified as SiO2 (EDS not shown).  Figure 10 ZSTS outer scale removed after exposure in stagnant air at 1927  C for 5 ten-min cycles. Left: outer surface of scale. Dark phase is SiO2 , bright phase ZrO2 and Zr-Ta-O. Right: underside of scale showing less SiO2 .     Ta2O5 into the cubic zirconia structure to form phase V results in a distorted orthorhombic structure [9] with all lattice parameters close to those of cubic zirconia. Clearly, Ta2O5 and possible excessive amounts of the phase V should be avoided. If Ta2O5 is present only in solution with ZrO2 , this would be avoided. The extent of Ta2O5 solubility in ZrO2 is unknown, but some +5 and Zr +4 ions are of solubility is likely since the Ta similar size (0.73 vs. 0.80 angstroms). A phase diagram for the related system ZrO2 -Nb2O5 system [10] shows 5 to 10 mol% solubility of Nb2O5 in ZrO2 . Currently, compositions with 5 vol% TaSi2 are being fabricated in hopes of avoiding liquid phase formation at ultra-high temperatures. Oxidation of this new material at 1927 C is planned.     4.3. Are TaSi2 additions the best way to add Ta to ZrB2 -based UHTC?  The ZSTC (TaC additions to ZrB2 -SiC) composition was formulated specifically to address this issue. Other possible vehicles for Ta additions include TaB2 and Ta5Si3 . A composition with 20 v/o TaC addition was chosen to see if the Ta additions offered improvements without additional B2O3 or SiO2 glass formation. As previously mentioned, the oxidation kinetics for ZSTC at 1627 C were rapid and linear indicating non-protective oxide formation, presumably due to scale porosity formed by evolution of CO and/or CO2 during the oxidation process. Results were not only much worse than the ZSTS composition, but also worse than the baseline ZS material. This is consistent     5975  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 11 ZSTS after testing in an arc jet at approximately 1890     evidence of melting.  C for ten min. Left: macro photos of 2.5 cm button. Right: surface oxide showing        with previous results obtained for the HfB2 -HfC system at 1500 C in air [11]. Protective scales were formed on HfB2 at 1500 C. The scales became less protective as HfC was substituted for HfB2 , with the most rapid oxidation for pure HfC. However, at the ultra-high temperatures achieved in arc jet testing, the HfC performed better than HfB2 . The better ultra-high temperature behavior of HfC was attributed to the tendency for the gaseous oxidation products CO and/or CO2 to be less disruptive to the HfO2 scale than boiling B2O3 . Oxidation of the ZSTC composition at 1927 C is planned to further evaluate this concept.     4.4. Can TaSi2 additions improve the oxidation resistance of HfB2 -based UHTC which are already superior to ZrB2 -based materials?  Oxidation resistance of the HS material was not improved, but actually worsened, by 20 v/o addition of TaSi2 . The HfO2 -Ta2O5 system appears to be even less studied than the corresponding ZrO2 system. However, there is evidence that the systems are similar in several phases, as well as formation of a Ta2O5 ·6HfO2 phase, respects. Some solid solution of the two oxide were both observed [12]. However, differences in oxidation rates of Hf metal and Zr metal have been documented. The parabolic oxidation rate constant of Hf metal has the same activation energy as that of Zr metal, indicating the transport of oxidant through HfO2 occurs by the same mechanism as that in ZrO2 ; however, the pre-exponential is lower indicating a lower number of defects available for oxygen transport in HfO2 [13]. Ta-additions may be less effective in HfO2 than in ZrO2 due to this possible difference in oxygen vacancy concentration. Another difference in the Zr-based and Hf-based SiC-containing UHTC is the upper temperature stability limit of ZrSiO4 versus HfSiO4 . The decomposition temperature of ZrSiO4 is not particularly well known.  5976                 ZrSiO4 decomposes to form ZrO2 and SiO2 at temperatures reported as low as 1538 C [14] or as high as 1676 C [15]. The XRD results from this study showed the absence of the ZrSiO4 phase for ZSTS after oxidation at 1627 C, consistent with the lower value of the ZrSiO4 stability temperature limit. The upper stability temperature of HfSiO4 appears to be higher than for ZrSiO4 although it is even less well studied. The only phase diagram found indicates HfSiO4 melts incongruently at 1750 C [16]. HfSiO4 was observed by XRD on HSTS after oxidation at 1627 C. The presence or absence of this MSiO4 phase may influence the oxidation kinetics. The oxygen transport rate in these MSiO4 phases is expected to be lower than in the respective MO2 oxides [17]. The oxidation rate for the HSTS material, which formed HfSiO4 , was higher than that of the ZSTS, which did not form ZrSiO4 . These results are inconsistent with a mechanism of reduced oxygen transport due to formation of an MSiO4 phase. The oxidized HSTS has not yet been analyzed microstructurally. Any differences in the oxide morphology and composition between the oxidized HS, HSTS and ZSTS materials will allow a better understanding of the effects of Ta-additions to the HS system. This microstructural work is also planned.  4.5. Can anything be learned about the mechanism by which Ta-additions improve the oxidation resistance of ZS materials at 1627 C?     More work is needed to answer this question. While glass immiscibility was observed for the ZSTS material after oxidation at 1627 C, it was not clear that tantalum was present in either of the phase-separated glasses. The phase identity and composition of the crystalline oxidation products remains uncertain due to energy overlaps in EDS Ta and Si peaks and similarity of XRD patterns for ZrO2 (c) and Ta2O5 ·6ZrO2 . Electron microprobe and/or WDS are needed to better characterize the oxidation products.  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  5. Summary and conclusions  The addition of 20 v/o TaSi2 to ZrB2 -SiC clearly improves the oxidation resistance of this material at 1627 C in air. The improved behavior is attributed to effects of Ta, not Si, although the mechanism for this mation due to melting of either Ta2O5 or Ta2O5 ·6ZrO2 improvement is not yet understood. Liquid phase forat 1927 C is a problem for the material with 20 v/o TaSi2 in ZrB2 -SiC. TaC additions are not effective in improving the oxidation resistance of ZrB2 -SiC at 1627 C. The addition of TaSi2 to the HfB2 -SiC system does not increase the oxidation resistance of this material at 1627 C in air. Future work includes the following: Measure the recession due to oxidation of UHTC materials to better quantify their oxidation rates. Characterize the oxidation products using WDS to better understand the amounts and compositions of phases present. Fabricate, oxidize and analyze UHTC materials with 5 v/o TaSi2 additions in hopes of avoiding liquid phase formation at temperatures greater than the melting point of Ta2O5 and Ta2O5 ·6ZrO2 . Continue to examine other compounds as vehicles for Ta-additions to UHTC.              Acknowledgments  The authors thank Don Ellerby (NASA Ames Research Center) and Matthew J. Gasch (Eloret, NASA Ames Research Center) for the arc jet testing and Ralph Garlick (NASA Glenn Research Center) for the XRD analysis.  References  1.  S . R . L E V I N E and E .  J . O P I L A , “Tantalum Addition to Zirconium Diboride for Improved Oxidation Resistance,” NASA/TM—  2003-212483 (2003).  2.  I . S .  G .  T A L M Y ,  J .  A .  Z A Y K O S K I , M . M .  O P E K A and  D A L L E K , “Oxidation of ZrB2 Ceramics Modified with SiC and Group IV-VI Transition Metal Borides,” in High Temperature  Corrosion and Materials Chemistry III, edited by M. McNallan and  E. Opila (The Electrochemical Society, Inc., Pennington, NJ, 2001)  p. 144.  3. M .  C A I L L E T ,  C .  D E P O R T E S ,  G .  R O B E R T ,  G .  V A L L I E R and G . V I T T E R , Rev. Int. Haute Temper. et Refract. 5  (1968) 173.  4. A . G . K O T L Y A R , A . D . N E U I M I N ,  S .  F .  P A L ’ G U E V , Z U B A N K O V ,  V .  N .  S T R E K A L O V S K I I  and V .  N .  Inorg.  Mater. 6(2) (1970) 281.  5.  S .  R .  L E V I N E , S I N G H , L O R I N C Z , P E T K O ,  M .  E .  J .  O P I L A ,  J .  A . J .  J .  D .  T .  E L L E R B Y  and M . G A S C H , “Ultra-High Temperature Ceramic Composites for Lead ing Edges,” JANNAF 39th Combustion/27th Airbreathing Propul sion/21st Propulsion Systems Hazards/3rd Modeling and Simulation  Joint Subcommittee Meeting, Colorado Springs, Dec. 1-5, 2003. 6. E . M . “Phase Diagrams for Ceramists” (The American Ceramic Society,  L E V I N ,  C .  R .  R O B B I N S and H .  F . M C M U R D I E ,  Inc., Columbus, OH, 1964) p. 144. 7. A .  R E I S M A N ,  F .  H O L T Z B E R G , M .  B E R K E N B I L T and  M . B E R R Y , J. Amer. Chem. Soc. 78 (1956) 4514.  8. B . W . K I N G ,  J .  S C H U L T Z ,  E . A . D U R B I N and W . H .  D U C K W O R T H , “Some Properties of Tantala Systems,” Battelle  Memorial Institute Report No. BMI-1106, 1956.  9. Powder Diffraction File, International Centre for Diffraction Data,  Swarthmore, Pennsylvania, card no. 42-60.  10.  “Phase Diagrams for Zirconium and Zirconia Systems”, edited by  H. M. Ondik and H. F. McMurdie (The American Ceramic Society,  Westerville, OH, 1998) p. 144. 11. A .  G .  M E T C A L F E , E . W U C H I N A , M .  N .  B .  E L S N E R , O P E K A and E .  D .  T .  A L L E N ,  O P I L A , “Oxidation of  Hafnium Boride,” in “High Temperature Corrosion and Materials  Chemistry,” edited by M. McNallan, E. Opila, T. Maruyama and  T. Narita (The Electrochemical Society Inc., Pennington, NJ, 2000)  p. 489. 12. R . 8(3) (2001) 560. P . K O F S T A D , “High Temperature Corrosion” (Elsevier Science Publishing Company, New York, 1988) p. 307. 14. C . 36(6) (1953) 190. 15. W . 52(5/6) (1967) 880. 16. E . M . F . M cM U R D I E , “Phase Diagrams for Ceramists, 1975 Supplement” (The American Ceramic Society, Inc.,  L . M A G U N O V and I .  R . M A G U N O V , Functional Mater.  13.  E .  C U R T I S and H .  G .  S O W M A N , J. Amer. Ceram. Soc.  C .  B U T T E R M A N and W .  R .  F O S T E R , Amer. Mineral.  L E V I N and H .  Columbus, OH, 1975) p. 165. 17. E . and E . Ceramic Oxides in the Range 1200  L . C O U R T R I G H T ,  J .  T .  P R A T E R , C . H . H E N A G E R  N .  G R E E N W E L L , “Oxygen Permeability for Selected     C-1700     C,” WL-TR-91-4006  (1991).  Received 30 October 2003 and accepted 31 March 2004  5977  \\x0c']"
},{
  "_id": 180,
  "PDF": "Oxidation of ZrB2-Based Ceramics in Dry Air.pdf",
  "Text": "['Oxidation of ZrB 2  -Based Ceramics in Dry Air   F. Monteverde and A. Bellosi   J. Electrochem. Soc.(cid:160) 2003, Volume 150, Issue 11, Pages B552-B559. doi: 10.1149/1.1618226  Email alerting service  Receive free email alerts when new articles cite this article sign up in the box at the top right corner of the article or  click here (cid:160)   To subscribe to  Journal of The Electrochemical Society  go to:   http://jes.ecsdl.org/subscriptions  © 2003 ECS The Electrochemical Society  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  (cid:160) (cid:160) \\x0c', 'B552  Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  0013-4651/2003/150~11!/B552/8/$7.00 © The Electrochemical Society, Inc.  Oxidation of ZrB2 -Based Ceramics in Dry Air  F. Monteverdez and A. Bellosi  National Research Council, Institute of Science and Technology for Ceramics, Faenza (Ra), Italy  The oxidation behavior of three different ZrB2 -based ceramics, a monolithic ZrB2 and two ZrB2 SiC composites, was studied up to 1350°C and under isothermal conditions at 1120°C for 20 h. Oxidation kinetics and microstructural changes in the oxidized specimens indicate that the monolithic ZrB2 has low thermal stability. On the contrary, the introduction of SiC particles markedly improves oxidation resistance due to the formation of an adherent and protective borosilicate glass layer that coats the sample surface, effectively limiting the inward diffusion of oxygen toward the reaction interface. The influence of the secondary grain boundary phases on oxidation is discussed. © 2003 The Electrochemical Society. @DOI: 10.1149/1.1618226# All rights reserved.  Manuscript submitted December 17, 2002; revised manuscript received May 26, 2003. Available electronically October 2, 2003.  ~i.e.,  Zirconium diboride ( ZrB2 ) is an emerging engineering ceramic because of its excellent and unique combination of properties. These include a high melting point, high elastic modulus, high electrical and thermal conductivity, good thermal shock and wear resistance, and good chemical inertness against acids, nonbasic slags, and nonferrous ~hot! metals. Such properties make it a candidate for numerous applications in high-temperature environments with aggressive gases and possible corrosive deposits.1 Moreover, ZrB2 and, by extension, borides of Hf and Ta, could be used for ultrahigh structural applications in aerospace, for example, thermal protection systems ~TPSs! for space vehicles in high oxygen potential environments.2-5 In addition to severe service conditions such as high temperatures and strong convective fluxes, TPS protections must be able to withstand intense mechanical stresses over the entire duration of the space mission vibrations at launch, landing impact!. Based on these introductory remarks, there is little doubt that TPS-dedicated materials, designed to be subjected to unusually hostile environments, demand high levels of dimensional and thermal stability. The remarkable physical stability of ZrB2 invariably requires the employment of pressure-assisted sintering techniques at very high temperatures ~.1900°C! to obtain highly dense components.1,2 Often the addition of sintering aids has been used to overcome the intrinsically low sinterability of highly refractory compounds. Highdensity materials were obtained through liquid phase sintering at processing temperatures well below those needed for undoped ZrB2 , and enhanced properties such as hardness, fracture toughness, and strength were also successfully achieved.6,7 The use of a ceramic-type additive ~i.e., silicon nitride! in a ZrB2 -based material showed overall improvement in mechanical properties when compared to literature on similar materials.1 In fact, the mechanical and physical properties of monolithic zirconium diboride are sometimes inappropriate for specific applications. The introduction of second phases ~i.e., SiC, ZrC! has succeeded in obtaining structurally stronger ZrB2 -based composites with improved resistance to ablation/oxidation.3-5,8-12 Complexshaped components of these materials are usually manufactured by electrical discharge machining. Alternative fabrication technology was recently investigated to produce near net shape components.13,14 As this class of advanced ceramics is designed to be exposed frequently to very hot and hostile environments, high-temperature stability, especially resistance to oxidation, is a key criterion for actual applications. The upper limit of the service temperature is closely related to the characteristics of grain boundary phases, which depend on the sintering aids and second phases that were used. For example, pure ZrB2 actively oxidizes above 1200°C due to intensive volatilization of whereas the SiC-containing ZrB2 showed stable oxidation behavior up to 1500°C. For temperatures  B2O3 , 15-17  z E-mail: fmonte@istec.cnr.it  above 1200°C, the addition of SiC provides more efficient oxidation resistance by encouraging the formation of borosilicate glass on exposed surfaces, which provided much more oxidation protection than B2O3 alone.3-5,9,11,14,18,19 Chemical interactions between the reaction scale ~formed during oxidation! and phases present in the bulk are very complex and strongly influence oxidation behavior. The growth of large bubbles underneath the external reaction scale often causes local spalling which, in turn, noticeably abates the mechanical properties. Concurrent release of gases ( B2O3 , CO) may lead to very low mass change, even when thick oxide scales form. Thus, comparison of the oxidation rate on a weight change basis must be accompanied by a microstructural examination of the oxide scale.5,18 This study focuses on the thermal stability of three hot-pressed ZrB2 -based materials, a monolithic and two ZrB2 SiC composites. Two types of thermal treatments in dry air were carried out in order to test the response of the three hot-pressed materials to oxidation, the first up to 1350°C and the second at 1120°C for 20 h. Oxidation kinetics was studied by measuring the weight change according to temperature and time. Microstructural modifications after oxidation were examined and correlated to the oxidation mechanisms involved.  Experimental  Fabrication  of  the  hot-pressed materials.—Starting  from the three highly dense ZrB2 -based mate compositions listed in Table I, rials were fabricated. The powder mixtures were homogenized for 24 h in a polyethylene bottle using ethanol and silicon nitride milling media, dried by rotating evaporator, and sieved. The as-processed powder batches were hot-pressed under vacuum ~;0.5 mbar! using BN-lined graphite die. Some properties of the hot-pressed materials are summarized in Table II. Detailed descriptions of the powder processing, microstructure, and mechanical properties of the three materials are reported elsewhere.6,10  Oxidation tests.—In gaseous environments, the widely accepted approach to quantify the response to oxidation is monitoring the weight change of a specimen over a certain duration at a given temperature.20 The fabricated materials were subjected to two ther~;50 cm3/min! at ambient presmal treatments in flowing dry air ~i! a nonisothermal sure, namely, run up to 1350°C, heating rate of 2°C/min, and free cooling; and ( an isothermal run at 1120°C ~heating rate of 20°C/min ! for 20 h and free cooling using a ther~model STA409, NETZSCH Gera¨ tebau mogravimetric analyzer GmbH, Germany!. The fast heating rate stage prior to the isothermal period was applied to minimize the oxidation effect before setting the target temperature of 1120°C. Test pieces with dimensions of 2.5 3 2.0 3 10.0 mm3 ~surface finish R a ’ 0.2) were cut out from the sintered billet, washed in an ultrasonic bath of acetone, dried at 80°C overnight, and finally placed inside an Al 2O3 vertical heated chamber on thin platinum wires. The mass variation was recorded  i i)  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  B553  Table I. Composition of the starting powder mixtures and hot pressing conditions: temperature T, holding time t, applied pressure P.  Material  Composition ~vol %!  T  ~°C!  t  ~min!  P ~MPa!  ZrB2 H.C. Starck grade B!  b-SiC ~H.C. Starck k, BF12!  a-Si3N4 ~Bayer, Baysinid!  Al2O3 ~Baikowski, CR30F Baikalox!  Y2O3 ~H.C. Starck k, grade C!  A B C  95 75 76.3  20 18.5  5 5 3.7  1  0.5  1700 1870 1760  10 15 15  30 30 30  ‘‘on-line’’ ( 1023 mg of accuracy! over the scheduled thermal run. The use of platinum as inert separator ensured minimal undesired reactions ~i.e., sticking effects! with the Al2O3 holder.  Analysis of  the microstructure.—Crystalline phases on the oxidized materials were analyzed using X-ray diffraction ~XRD; model D500 Siemens, Germany! with a Ni filtered Cu Ka radiation at 25 mA and 30 kV. Polished surfaces ~finished 1 mm! of both sintered and oxidized specimens were first prepared using diamond abrasives and nonaqueous lubricant, then observed via scanning electron mi~SEM; U.K.! croscopy model S360, Cambridge Instruments, ~EDX; coupled with an energy-dispersive X-ray microanalyzer model INCA Energy 300, Oxford Instruments, U.K.!. The crosssectioned ~polished! surfaces of the oxidized coupons, imaged using secondary ~SE! and backscattered ~BSE! electron signal, were observed free of conductive coating in order to maintain the sensitivity of the EDX equipment to the low Z elements as high as possible. In fact, detecting light elements helps trace the diffusion paths of the moving species and may clarify the chemical reactions controlling the oxidation kinetics. Some thermodynamic calculations were done using the software package HSC 5.0.21  Results  Microstructure of hot-pressed materials.—The final density of  material A is 5.85 g/cm3. This corresponds to a relative density of 98% in accordance with the mixture rule. The microstructure is fairly regular, with low residual porosity. Besides ZrB2 , XRD-SEMEDX analyses found a limited amount of grain boundary phases: crystalline BN, tetragonal ZrO2 , and a glassy compound in the system B-N-O-Zr-Si ~Fig. 1!. These secondary phases formed during hot-pressing derive from reactions among Si3N4 ~used as a sintering aid!, ZrB2 , boron and silicon oxide ~present on ZrB2 and Si3N4 particles, respectively!. As with TiB2 doped with Si3N4 , 22 the liquid phase formed during the hot-pressing of material A is supposed to act as medium for partial dissolution of the starting ZrB2 particles, volume diffusion of the involved species, and reprecipitation on undissolved ZrB2 particles.6 Then, cooling down from the hot-pressing temperature, the mentioned grain boundary phases remain mainly at triple points or in randomly distributed pockets. With regard to composite B and C, the need for a higher hotpressing temperature to achieve near full density ~Table II! was primarily due to the dragging action of the SiC particulate on the grain boundary mobility during sintering. The introduction of extra addi tives such as Al2O3 and Y2O3 in material C visibly diminished the melting point of the liquid phase and allowed us to obtain a dense ZrB2 SiC composite at a lower hot-pressing temperature, if compared to material B.10 The final textures of both composites are fairly similar ~Fig. 2a and b!, with grain size of ZrB2 grains on average smaller than that of material A. The residual porosity is very low, while the majority of darker regions in the SEM images ~Fig. 2! consists of BN or BN-rich SE phases. The SiC particles generally exhibit a clustered distribution and are often in contact with reaction products of the sintering process. More microstructural details are reported elsewhere.10  Weight change during oxidation tests.—Specific weight change  vs.  T max 4 Tmin  temperature of the three tested materials during the nonisothermal run up to 1350°C is displayed in Fig. 3. For comparison, weight change data of hot-pressed pure ZrB2 23 are presented in the same graph. The heating rate of 2°C/min allowed us to measure not only the temperature at which the sample begins to gain weight ( T ON) but also the temperature corresponding to the change of the slope in the thermogravimetric ~TG! curve ( T MAX , T MIN) . An inset in Fig. 3 shows the values of T ON , T MIN , and T MAX . The observed local variation of the slope in the TG curves is most likely due to the thermal instability of the oxide scale growing on the external faces of the oxidizing sample. Such behavior has been reported by other authors on similar systems.24 Therefore, the temperature of the isothermal run, i.e., 1120°C, selected within the range, corresponds to oxidizing conditions in a partial protective regime for materials being tested. The net weight change during 20 h of exposure at 1120°C is shown in Fig. 4. The initial offset of the specific weight gain accounts for oxidation during the early heating stage. The TG data show dissimilar responses to oxidation for the three materials: the weight gain of the monolithic ZrB2 is greater than that of the ZrB2 SiC composites ~B and C!. Material B did not experience appreciable net mass change after 20 h of exposure. Some important results were obtained when applying the multiple-law model developed by Nickel25 to the analysis of the weight gain data. For material A the fit of the TG values conferring the most realistic solution ( R 2 5 98.6) is displayed in Fig. 5 and accounts for paralinear kinetics. An essential output of this calculation stands in the negative linear contribution over the entire duration of the test. From 500 to 800 min of exposure, neither mass gain nor mass loss reactions predominate ~Fig. 4!, mainly because the  ~A!  linear thermal expansion coefficient l (cid:132)25-1300°C(cid:133), Table II. Properties of the hot-pressed materials: mean grain size d, electrical resistivity r, Vickers microhardness Hv 1.0, fracture toughness KIc, 4 pt flexural strength in air at different temperatures.  Sample  Density  d  r  l  Hv1.0  ~g cm23!  5.85 5.30 5.32  ~%  98 98 98  !  m  ~m  !  ~mVcm  !  1026/°C  ~  !  GPa!  ~  3 2.4 2.5  7 15 16  7.43 7.50  13.4 6 0.6 14.6 6 0.3 14.2 6 0.6  A B C  KIC  MPa Am  3.75 6 0.10 4.55 6 0.10  s  25 °C  1000°C  1200 °C  600 6 90 730 6 100 710 6 110  400 6 20 430 6 110 630 6 110  240 6 30 250 6 10 280 6 20  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'B554  Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  Figure  ZrB2 ,  1. SE-SEM micrograph from a polished ~2! ZrO2 , and ~3! BN or B-N-Si-O-Zr.6  area  of material A:  ~1!  growth of the oxide scale depends to a large extent on the exposure time. A parallel isothermal run performed in flowing pure oxygen at 1100 °C provided a monotonically decelerating trend of the TG data ~Fig. 6!. The excellent fitting solution ( R 2 5 99.8) confirmed that in such a temperature range, oxidation behavior can be described properly using a paralinear model. In the case of composite B, after a transitory period of 400 min, resistance to oxidation ~in terms of weight change! resulted in near insensitivity to the oxidation attack ~Fig. 4!. However, material C showed a prevailing linear dependence of the TG data vs. exposure time ~Fig. 4!.  Figure 2. SE-SEM micrographs from a polished area of ~a! material B and ~b! C: ~1! ZrB2 , ~2! ZrO2 , ~3! SiC, and ~4! BN-based secondary phases.10  Microstructure  tests. — Nonisothermal  after  oxidation  run.—After the nonisothermal test, extensive microstructural modifications occurred in material A. XRD analysis detected highly textured monoclinic ZrO2 phase on the exposed surfaces. Moreover, the characteristic hump in the XRD pattern denoted the existence of a glass-like reaction product. Examination of the cross-sectioned surface of the oxidized piece showed an irregular oxide scale, characterized by large cavities, bubbles, glass infiltration, and evidence of spallation ~Fig. 7!. Closer inspection of the external oxide scale via SEM-EDX revealed the presence of ZrO2 particles enclosed by borosilicate glass ~Fig. 8!. Underneath such a layer, internal bubbles are sign of gas evolution during thermal treatment. In comparison to ~monolithic! material A, the addition of silicon carbide markedly improved resistance to oxidation of the ZrB2 SiC composites ~B and C!, as verified by much thinner oxide scales ~Fig. 9!. XRD analysis on the oxidized surface detected tetragonal ZrO2 and ZrSiO4 in addition to strongly oriented monoclinic ZrO2 crystals ~Fig. 10!. XRD peak intensities of the former phase are more pronounced in material C than in material B ~explanation is given in the discussion section!. With reference to the reaction scale arrangement of material B and C, a multilayered assemblage was observed. The basic scheme for both composites consists of an adherent outermost glass layer covering a ZrO2 -based interlayer embedded in a glass melt ~Fig. 11!. The EDX study of the glass basically assessed the composition of a borosilicate. Discontinuous infiltration or pockets of glass were observed, whereas no evidence of bubbles was found.  Isothermal run.—The magnitude of the weight change data ~Fig. 4! is consistent with the microstructural modifications that the oxidized materials underwent during the isothermal test at 1120°C for 20 h of exposure.  Figure 3. Specific weight change ~w! vs. temperature ~T! of the three tested materials. Data of a pure ZrB2 material were plotted for comparison.  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  B555  Figure 4. Specific weight change ~w! vs. time of exposed for 20 h at 1120°C in flowing dry air.  the three tested materials,  In material A, a rather compact scale consists of ZrO2 particles covered with a glassy film ~Fig. 12!. An intermediate layer, characterized by ongoing traces of oxidation, separates such a ZrO2 -based region from the unreacted bulk. A ZrO2 -free thin glass layer survives on the top of the reaction scale. The reaction scales of composite B and C look very similar ~Fig. 13!. A coherent glass layer covers the inner portions of the bulk. An intermediate layer, much thinner than that in material A, shows clear traces of oxidation gradually attacking deeper regions of the unreacted material. Surprisingly, several SiC particles lying close to the outermost glassy layer seem to be unaffected by the effects of the oxidation attack. Among the oxidation products, highlighted in detail in Fig. 14, monoclinic and tetragonal ZrO2 ~less oriented than after the nonisothermal treatment! represent the main crystalline new formed phase in all the studied materials. A minor content of ZrSiO4 was detected, but only in material B and C. The glass basically consists of a borosilicate. Its main characteristics, such as composition, viscosity, and vapor pressure, are expected to change according to time. The amount of boron in such a melt represents a parameter that greatly  Figure 6. Specific weight change ~w! vs. exposure time ~t! of material A at 1100 °C in flowing oxygen. Single contribution from the mixed law w5A1K PARA t 1 K LINt , A constant, are plotted.  affects its physical properties. The EDX elemental analysis of the glass in material B and C, apart from boron, silicon, and oxygen, stated the presence of C and of Al 1 C, respectively.  Discussion  The results described show that environmental stability, namely, oxidation resistance of ZrB2 -based ceramics in flowing dry air, depends on the microstructural characteristics of the as-sintered materials. Specifically, the addition of SiC particles is the key factor that makes the tested ZrB2 SiC composites ~B and C! consistently more resistant to oxidation than the monolithic ZrB2 . Taking the phases present in the as-sintered materials into consideration, both condensed and gaseous oxidation products can form. The expected main reactions describing the oxidation process, which involve either mass gain  ZrB2 1  5  2  O2 5 ZrO2 1 B2O3 ~ l!  2 BN 1  SiC 1  3  2  3  2  O2 5 B2O3 1 N2  O2 5 SiO2 1 CO  or mass loss  B2O3 ~ l! 5 B2O3 ~ g!  @1#  @2#  @3#  @4#  Figure 5. Specific weight change ~w! vs. exposure time ~t! of material A, exposed at 1120°C in flowing dry air for 20 h. Single contribution from the ~A constant!, are plotted.  mixed law w 5 A 1 K PARA t 1 K LINt  Figure 7. BSE-SEM micrograph from polished section of material A, after the nonisothermal run: ~1! oxide scale, ~2! cavities, bubbles, ~3! infiltration of glass, and ~4! unreacted bulk.  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'B556  Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  Figure 8.  ~Left! SE-SEM micrograph ~2 kV! from an outermost polished portion of the oxide scale in material A, after the nonisothermal run. ~Right! EDX spectra ~4 keV! from the numbered features in the micrograph.  SiO2 ~ s! 1 CO~g) 5 SiO~g) 1 CO2 ~ g!  @5#  are simultaneously active. Therefore great care was taken when interpreting weight change data due to the formation of volatile products. The main reaction products from the oxidation of ZrB2 , which constitutes the basic matrix of the tested materials, are zirconia and amorphous boric acid ( B2O3 ) . In fact, the latter compound has an 26 and a high vapor pressure unusually low melting point, 450°C, ~Fig. 15!, so at high temperature it quickly vaporizes. The rapid increase in weight gain of a pure ZrB2 material with 87% relative density ~Fig. 3! reflects the weak ability of B2O3 to hinder intensive oxidation. ZrO2 has a semiprotective action at elevated temperatures due to its anion-deficit structure permitting inward transport of oxyvacancies.20 Nonetheless, gen via lattice grain boundary siliconcontaining phases, although in very limited amounts as in material A, promote the formation of a borosilicate glass with higher melting point, and lower vapor pressure ~Fig. 15! and viscosity ~Fig. 16!. The global effect of this borosilicate glass is to improve resistance to oxidation by acting as a barrier against the inward diffusion of oxygen more efficiently than B2O3 and ZrO2 . The creation of the two well-defined layers, glass and glass plus ZrO2 particles, in composites B and C ~Fig. 7-9! suggests that the wetting angle between the borosilicate glass and ZrO2 is very high.3 This feature not only ensures an almost continuous coating of the sample but also a cov Figure 10. Unprocessed XRD patterns from the oxidized surfaces of material A and B, after the nonisothermal run up to 1350 °C. The ZB peak arises from the unreacted bulk.  erage of ZrO2 particles constituting the oxide subscale, both providing increased inhibition of the inward access of oxygen. The aforementioned microstructural modifications and paralinear kinetics are representative of the limited oxidation resistance of material A. The finding of a negative linear contribution ~Fig. 5 and 6! in the oxidation kinetics accounts for a release of volatile products, most likely the boria fraction from the outermost borosilicate layer. Additionally, the parabolic contribution can be explained in terms of the growth of the external oxide scale, which progressively imposes longer diffusion paths for oxygen to reach the reaction interface. The micrograph in Fig. 17 shows that the deepest front of the oxidation  Figure 9.  ~a! BSE-SEM micrographs from polished section of the external oxide scale of ~a! material B and ~b! C after the nonisothermal run. White lines demarcate the outermost glass layer, while the arrow indicates the internal infiltration of glass.  Figure 11. SE-SEM micrograph ~2 kV! from a polished area of the oxide scale in material B, after the nonisothermal run: ~1! outermost glassy layer, ~2! ZrO2 -based interlayer, and ~3! unreacted bulk.  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  B557  Figure 13. SE-SEM micrograph ~2 kV! from a polished section of the oxide scale of material C, after the isothermal treatment: ~top! outermost glassy layer and ~bottom! intermediate layer.  particles acts as an effective obstacle against the inward diffusion of oxygen along grain boundary channels and at the same time, helps form a silica-enriched glass. This melt, fluid at the testing temperature and characterized by a restricted permeability to oxygen, quickly seals the external surfaces of the test piece and short-circuit paths for the incoming oxygen ~i.e., residual porosity, cracks! toward easily accessible ZrB2 grains and/or grain boundary phases. The evidence of residual unoxidized SiC particles within the reaction scale settled the fundamental role of such a phase in suppressing the advance of the oxidation attack. In comparison with material B, the moderately lower resistance to oxidation of material C can be ascribed to the modified characteristics of the external glass layer. In fact, Y and Al cations contained in the secondary phases of material C diffuse out toward the reaction interface during oxidation tests, varying the properties of and the glass!. On one side, Y ions the oxidation products ( ZrO2 stabilize a tetragonal ZrO2 polytype and this can account for the higher peak intensity of t-ZrO2 in the XRD spectra of the oxidized surface in material C, compared to material B. On the other side, the enrichment in such ion modifiers causes the partial breakup of the silicon-oxygen network of the glass, and invariably alters properties such as viscosity and melting point, with an overall enhanced capa Figure 14. BSE-SEM micrograph from the oxide scale of material C, after the isothermal run: ~1! ZrB2 , ~2! ZrO2 , outermost glass ZrSiO4 , and ~5! SiC. The floating bright particles on the top of layer are Y-stabilized ZrO2 .  the glass  ~3!  layer,  ~4!  Figure 12. SE-SEM micrograph ~2 kV! from a polished section of the oxide scale of material A, after the isothermal treatment: ~1! ZrO2 -based layer, ~2! intermediate layer, and ~3! unreacted bulk. Black arrows on the top indicate the thin glass layer.  attack advances along grain boundaries toward the inner regions of the material. Grain boundaries act as preferential paths for the outward migration of boron and silicon ~the former more rapidly! and for the inward transport of oxygen, inducing a continuing conversion of ZrB2 into ZrO2 . Besides the evaporation of mass from the oxide scale, the concurrent action of other physical factors ~i.e., solution of gases in the the diffusing species! often scale, changing diffusion rate/path of leads the predicted kinetics model to depart from the experimental data. The atypical behavior of the isothermal TG curve in material A ~Fig. 5! can be explained by assuming the occurrence of a temporary decay of the oxide scale protectiveness, i.e., breakaway reaction, which derives from the partial spalling of the oxide, particularly in correspondence to the bursting of the gaseous bubbles ~Fig. 7! on the near side of the reaction interface. After a relatively long exposure time ~about 500 min!, the achievement of a critical amount/ thickness of the oxidation products makes the volatilization of gaseous products ~i.e., mass loss! the physical phenomenon dominating the resulting net weight change. Similar behavior was described on a TiB2 system, doped with Si3N4 as sintering aid.24 With regard to the resistance to oxidation of material B and C, the presence of silicon carbide as a particulate second phase is certainly responsible for the noticeable improvement in oxidation resistance. In the very first oxidation step at 1120°C, the majority of SiC  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'B558  Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  Figure 15. Vapor pressure vs. ambient pressure.21  temperature above some oxides, calculated at  bility of transporting oxygen into the bulk.27 The resultant linear trend of the TG data for material C supposedly accounts for the reaction at the boride/oxide interface as the rate-limiting step governing the response to oxidation at the tested temperatures.  Conclusions  Three highly dense ZrB2 -based materials, a monolithic and two SiC-containing composites, were fabricated by hot-pressing using silicon nitride as sintering aid. Resistance to oxidation, tested in flowing dry air up to 1350°C ~2°C/min! or for 20 h at 1120°C, was monitored measuring the weight change vs. temperature or vs. exposure time. In all the studied materials, the main oxidation products were ZrO2 and a borosilicate glass. The monolithic material ( ZrB2 1 5 vol % Si3N4 ) suffered microstructural modifications that indicated a low oxidation resistance. The external oxide scale partially protects the inner portions of the bulk. The experimental weight change data agree with a paralinear kinetics The introduction of SiC particles markedly improved the oxidation resistance of the ZrB2 SiC-based composites in comparison to the monolithic ZrB2 . An adherent and protective borosilicate glassy layer coats the sample surface, greatly limiting the inward diffusion of oxygen into the unreacted bulk. The addition of extra additives ( Al2O3 and Y2O3 ) in the SiCcontaining composite lowered the viscosity of the outermost glassy  Figure 16. Viscosity ~h! of two glass-forming compounds SiO2 and B2O3 as a function of reciprocal temperature.27  Figure 17. SE-SEM micrograph ~2 kV! from inner portions of the oxidation front in material A, after the isothermal run. The vertical arrow indicates the advancing direction of the oxidation attack.  scale and therefore its permeability to oxygen, in comparison with an undoped borosilicate glass. The overall effect was a reduced capability to resist the oxidation attack.  Acknowledgments  The authors thank Dr. G. Laudisio ~Institut fu¨ r Mineralogie, Petrologie und Geochimie, Eberhard Karls Universitaet, Tuebingen! for the tests in flowing oxygen, and G. Ercolani and D. Dalle Fabbriche ~ISTEC-CNR, Faenza! for technical assistance on the thermal treatments.  The National Research Council, Institute of Science and Technology for Ceramics, Italy, assisted in meeting the publication costs of this article.  References  1. C. Mroz, Am. Ceram. Soc. Bull., 74(cid:132)6(cid:133), 164 ~1995!. 2. K. Upadhya, J.-M. Yang, and W. Hoffman, Am. Ceram. Soc. Bull., 58 , 51~1997!. 3. M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey, J. Eur.  Ceram. Soc., 19, 2405 ~1999!.  4.  S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem, J.  Eur. Ceram. Soc., 22, 2757 ~2002!.  9. 10.  5. C. R. Wang, J.-M. Yang, and W. Hoffman, Mater. Chem. Phys., 74, 272 ~2002!. 6. F. Monteverde and A. Bellosi, Scr. Mater., 46, 223 ~2002!. F. Monteverde, A. Bellosi, and S. Guicciardi, J. Eur. Ceram. Soc., 22, 279 ~2002!. 7. 8. I. G. Talmy, J. A. Zaykoski, and M. A. Opeka, Ceram. Eng. Sci. Proc., 19(cid:132)3(cid:133), 105 ~1998!. J. Bull, M. J. White, and L. Kaufman, U.S. Pat. 5,750,450 ~1998!. F. Monteverde, S. Guicciardi, and A. Bellosi, Mater. Sci. Eng., A, 346(cid:132)1-2(cid:133), 310 ~2003!. 11. I. Ogawa and T. Yamamoto, J. Mater. Sci. Lett., 11, 296 ~1992!. 12. G.-J. Zhang, Z.-Y. Deng, N. Kondo, J.-F. Yang, and T. Ohji, J. Am. Ceram. Soc., 83, 2330 ~2000!. 13. N. B. Dahotre, P. Kadolkar, and S. Shah, Surf. Interface Anal., 31, 659 ~2001!. 14. C. Bartuli, T. Valente, and M. Tului, Surf. Coat. Technol., 155, 260 ~2002!. 15. W. C. Tripp and H. C. Graham, J. Electrochem. Soc., 118, 1195 ~1971!. J. B. Berkowitz-Mattuck, J. Electrochem. Soc., 113, 908 ~1966!. 16. 17. M. Singh and H. Wiedermeier, J. Am. Ceram. Soc., 74, 724 ~1991!. 18. W. C. Tripp, H. H. Davis, and H. C. Graham, Ceram. Bull., 52, 612 ~1973!.  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c', 'Journal of The Electrochemical Society, 150 ~11! B552-B559 ~2003!  B559  19. K. Kobayashi, H. Sano, K. Maeda, and Y. Uchiyama, J. Ceram. Soc. Jpn., 100, 398 ~1992!. 20. K. Bundschuh and M. Schutze, Mater. Corros., 52, 204 ~2001!. 21. HSC Chemistry for Windows 5, Outokumpu Research Oy, Pori, Finland. L.-H. Li, H.-E. Kim, and E. S. Kang, J. Eur. Ceram. Soc., 22, 973 ~2002!. 22. 23. A. Bellosi, F. Monteverde, D. Dalle Fabbriche, and C. Melandri, J. Mater. Process.  Manuf. Sci., 9, 156 ~2000!.  24. Y.-H. Koh, S.-Y. Lee, and H.-E. Kim, J. Am. Ceram. Soc., 84, 239 ~2001!.  25. K. G. Nickel, in Corrosion of Advanced Ceramics/Measurement and Modelling, K.  G. Nickel, Ed., p. 59, Kluwer Academic Publishers, Norwell, MA ~1994!. 26. M. W. Chase, Jr., C. A. Davis, J. R. Downey, D. J. Frurip, R. A. MacDonald, and A. N. Syverud, JANAF Thermochemical Tables, 3rd ed., American Chemical Society and American Physical Society, New York ~1986!. 27. R. H. Doremus, in Glass Science, p. 105, John Wiley & Sons, New York ~1973!.  Downloaded on 2012-09-25 to IP   128.143.23.241   address. Redistribution subject to ECS license or copyright; see   www.esltbd.org  \\x0c']"
},{
  "_id": 181,
  "PDF": "Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 32 (2012) 3875-3883  Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content  Peter A. Williams a , Ridwan Sakidja, John H. Perepezko, Patrick Ritt  ∗  Department of Materials Science & Engineering, University of Wisconsin-Madison, Madison, WI 53706, United States  Received 25 January 2012; received in revised form 14 May 2012; accepted 15 May 2012  Available online 22 June 2012  Abstract  The oxidation performance of ZrB2 -SiC ultra-high temperature ceramics with SiC content ranging from 20 to 80 vol% has been evaluated at 1773 K for 50 h and at 2073 K for 20 min. Oxidation reaction pathways were  interpreted using volatility diagrams of  the ZrB2 -SiC system. At 1773 K ≤50 vol% SiC developed a for 50 h, all ZrB2 -SiC composites from 20  to 80 vol% SiC formed a protective SiO2 surface coating. Samples with  distinguishable SiC-depleted layer at 1773 K and 2073 K. High temperature torch testing for 20 min at approximately 2073 K revealed that samples ≥65 vol% SiC exhibit a depression under  with  the  torch flame. Samples rich  in ZrB2 were dominated by a ZrO2 layer after a similar exposure. The overall weight density of ultra-high temperature ceramics can be reduced with improved oxidation performance at 1773 K by adding at least 65 vol% SiC. © 2012 Elsevier Ltd. All rights reserved.  Keywords: B. Microstructure-final; C. Thermal properties; D. Borides; D. SiC  1.   Introduction  The need  for advanced materials  in aerospace applications has been an  important driver of  innovation  in protective coatings and processing methods. The advent of hypersonic flight certainly qualifies as a huge step  forward  in performance and capability, and this represents a significant materials capability challenge. Reentry vehicles, regardless of their specific designs, require control surfaces with sharp leading edges if they are to be highly maneuverable at hypersonic velocities. Low-radius leading edges are subject to much greater aerothermal heating than blunt edges such as  those on  the Space Shuttle.1,2 In order  to limit the rate of diffusion controlled processes such as oxidation and creep during ultra-high temperature operation, the operating temperature should be 0.67 Tm or lower which, in many applicatemperature above 2923 K.3 Available tions,  implies a melting   ∗  Corresponding author at: 1213 Engineering Research Building, 1500 Engi neering Dr., Madison, WI 53706, United States. Tel.: +1 608 265 4986;  fax: +1 608 263 1660.  E-mail addresses: peter.a.williams@intel.com (P.A. Williams),  pritt@wisc.edu (P. Ritt). a Now employed by Intel Corporation, 5000 W. Chandler Blvd., Mail Stop  C5-165, Chandler, AZ 85226, United States.  0955-2219/$ - see front matter © 2012 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2012.05.021  thermal protection materials will not survive such extreme temperatures for extended service, and new materials are required for advanced thermal protection systems. Currently, the ultra-high temperature ceramic (UHTC) materials receiving  the most attention as possible  leading edges for hypersonic vehicles for use  in ultra-high  temperature environments are  the  refractory metal diborides, zirconium diboride (ZrB2 )  and  hafnium  diboride  (HfB2 ). Because  of  the  high cost and density of HfB2 as opposed  to ZrB2 (10.5 g/cm3 and 6.1 g/cm3 , respectively), ZrB2 has been favored as a candidate UHTC  in  recent studies.4-10 The properties of  the  refractory metal  diborides  such  as  high melting  point,  thermal  shock resistance, and high thermal conductivity have made them attractive  for hypersonic  applications.11 The primary  concern  for these materials, however, is the inability of the refractory metal diborides  to adequately  resist oxidation at ultra-high  temperatures. A  thorough model for  the oxidation of refractory metal diborides (including ZrB2 ) was presented by Parthasarthy et al.12 It was shown that for the oxidation of ZrB2 : ZrB2 +  ZrO2 +  B2O3 (l)  ≈1273 K,  At  temperatures above  the oxide scale becomes less protective because of the preferential evaporation of boron oxide (boria, B2O3 ). In  the 1373-1673 K  temperature regime,   5/2O2 (g)   ⇔  (1)        \\x0c', '3876   P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883  Fig. 1. Cross-sectional SEM images of as-sintered S20Z80 (a) and S80Z20 (b). Dark regions represent SiC and light regions represent ZrB2 .  ×  =  ρZrB2  the evaporation of B2O3 is  rapid which  leads  to  linear oxidation kinetics of  the ZrB2 substrate.13-17 Consequently, silicon carbide  (SiC)  is often added  to create  the ceramic composite: (refractory metal)B2-SiC. The addition of SiC  to  the  refractory metal diborides has been shown  to  improve  the oxidation resistance of  the UHTC due  to a continuous borosilica scale with a  lower oxygen diffusion rate  than boria18 that forms on the surface.14,19,20 For a 80 vol% ZrB2 -20 vol% SiC composition at 1823 K  the oxygen diffusivity  in  the borosilica has −14 m2 /s for a B2O3 -21 mol% SiO2 been evaluated as 1.7   10 composition which is 107 times higher than that for pure silica, in ZrO2 (10 −10 m2 /s at but much  lower  than  that for oxygen  1773 K).7 The refractory ZrB2-SiC system also has a eutectic melting temperature near 2573 K.21 To date, most of  the published  experimental  research on ZrB2-SiC UHTCs has  focused on a narrow  range of compositions on  the ZrB2 end of  the pseudo-binary (i.e. 10-30 vol% SiC).4-10,14,18,19,22-24 There are, however, apparent advantages to SiC-rich composites  that  should be addressed. For example, compositions rich  in SiC are significantly  less dense  than compositions on  the other end of  the ZrB2-SiC pseudo-binary (ρSiC = 3.21 g/cm3 and   6.12 g/cm3 ). Also,  additional SiC  in  the system may provide enhanced oxidation protection because the oxidation product of SiC, SiO2 , is the primary barrier  to oxidation for  the system. This  is only  true, however, up to  the point at which  the oxidation protection of SiC breaks 1923 K due to the active-passive transition, which down above  is  triggered by a reaction at  the SiC-SiO2 interface.25 Furthermore, previous literature indicates that significant evaporation of SiO2 (l) to SiO(g) can occur at around 2048 K under atmospheric pressure.18,26 In addition to the environmental stability at high temperature the development of a SiC-depleted  layer has been  frequently observed  for  the oxidation of ZrB2 -rich  compositions  ranging  from 10  a model analysis of  the volatility diagram for  the ZrB2 -SiC system Fahrenholtz attributed the formation of this region to a  gradient  that develops underneath  the surface during oxidation in air.4 However,  the development of  the SiC-depleted  region is not a universal phenomenon. For example, in HfB2-20 vol% SiC composites a SiC-depleted zone was not detected following arc jet testing.29 In this case it was suggested that the distribution of SiC particles may play an important role.  to 30 vol% SiC.4,6,8-10,14,24,27-29 Based upon   pO2  In  the current study,  the oxidation performance of UHTCs on both ends of  the ZrB2 -SiC pseudo-binary  is examined  in order to more closely identify optimal ZrB2 -SiC compositions in relation to the expected operating conditions up to 2073 K.  2. Materials and methods  2.1. UHTC synthesis  The UHTC samples were prepared using  individual SiC (␣ polymorph, hexagonal crystal structure) and ZrB2 (hexagonal crystal  structure) powders  from Alfa Aesar. The ZrB2 pow−325 mesh  (≤44  der was  \\u242em particle diameter) with 99.5% purity, and  the SiC powder was reported  to have a mean parti\\u242em with 99.8% purity. The volume of the weighed cle size of 2  powders was measured  for each of  the nominal compositions using a graduated cylinder then poured into a beaker with pure ethyl alcohol. The suspension was mixed using a magnetic stir plate then dried yielding a powder of each nominal composition. While some local composition variation may occur with this procedure due to incomplete agglomerate break up, examination of hot-pressed samples by scanning electron microscope  (SEM) JEOL JSM 6100 with a 4 nm resolution and energy dispersive X-ray spectroscopy (EDS) WINEDS with an energy resolution of 138 eV did not reveal significant agglomerates. Representative microstructures for  the high ZrB2 (S20Z80) and high SiC (S80Z20) samples are shown  in Fig. 1(a) and (b), respectively. Table 1 shows the compositions examined as well as the sample identification. For  all  samples  except  the  S80Z20  sample,  the mixed powder was then poured into a machined graphite die and compressed under 30 MPa using  a hydraulic hand press  (Model B, Carver  Inc., Wabash,  IN). The die was  then placed  into a  Table 1  UHTC compositions and their name used throughout the paper.  Composition   80 vol% SiC-20 vol% ZrB2 65 vol% SiC-35 vol% ZrB2 50 vol% SiC-50 vol% ZrB2 35 vol% SiC-65 vol% ZrB2 20 vol% SiC-80 vol% ZrB2  Name in this study  S80Z20  S65Z35  S50Z50  S35Z65  S20Z80  \\x0c', 'P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883   3877  10  SiO2()l  + ZrB2  )  x  p  (  g o  l  5  0  −5  −10  −15  −20  SiO2()l  + ZrO2 +  B2O3()l  SiC + ZrB2  SiO(g)  1500°C 1650°C 1800°C  −30  −20  B2O3(g)  BO2(g)  r  i  a  n  i  2  O  p  0  10  −10  log(pO2  )  Fig. 3. Volatility diagram   the ZrB2 -SiC system at 1500 C, 1650 × 105 Pa). 2073 K. The pressures are relative to ambient pressure (1.01   for         C, and  A disappearing filament pyrometer (The Pyrometer Instrument ±0.5% was used Company, Windsor, NJ) with an accuracy of  during the tests to measure the sample temperature. The gradient experiments were conducted at 2073 K under the cone of the flame with the edges of the sample reaching about 1873 K. The cone diameter was approximately 1.5 cm, and the distance from the  torch opening  to  the sample was approximately 3 cm. The sample size was the same as that used in the isothermal oxidation testing. UHTC samples were exposed to the torch flame for 20 min  then air cooled. The  tests were duplicated on separate samples of  the  same nominal composition. Like  the  isothermal tests, all of the oxidized samples were cross-sectioned and polished for characterization in the SEM.  3. Calculation  3.1. Volatility diagram  A volatility diagram for the ZrB2-SiC system (Fig. 3) was calculated using the vapor pressures of the relevant gaseous species (Table 2). The diagram was compiled using data from the JANAF tables30 and HSC 6.1 software31 and  is based on  the work of Heuer and Lou32,33 and Fahrenholtz.4 It was assumed that condensed phases do not react with each other and, therefore, have unit activity. This assumption was also made for  the oxidation products  (i.e. SiO2 , ZrO2 , and B2O3 ). The same assumptions were used by Fahrenholtz.4  4. Results  4.1.   Isothermal furnace study  All five of  the samples  in Table 1 were oxidized at 1773 K for 50 h in static air. The Archimedes method was used to measure  the density of  the hot-pressed and spark plasma sintered billets that ranged from 90% to 94% of the theoretical densities based upon  the nominal compositions. The density variations  Fig. 2. Schematic of the set up used for torch testing in this study.  tungsten-mesh  element  hot-press  and  uniaxially  pressed  at −5 mbar). The 7 MPa for 10 h at 1973 K under high vacuum (<10 hot-pressed cylindrical billets (2.54 cm diameter) were cut using a diamond saw (Model VC-50, LECO Corp., St. Joseph, MI) to 0.3 cm thick coupons. The coupons were ultrasonically cleaned in pure ethyl alcohol then dried. The S80Z20 powder was spark-plasma sintered (FCT Systeme HP D25, Rauenstein, Germany) for 5 min at 2373 K under a pressure of 50 MPa using a ramp rate of 200 K/min at Alfred University  (Alfred, NY). The spark-plasma sintered cylindrical billets  (2.06 cm diameter) were examined at  their original thickness of 0.55 cm. The S80Z20 coupons were ultrasonically cleaned in pure ethyl alcohol then dried.  2.2.   Isothermal oxidation  The isothermal oxidation tests were done in a MoSi2 element furnace with an open alumina  tube (Deltech Inc., Denver, CO) capable of reaching 1973 K. For this study, isothermal tests were performed at 1773 K in stagnant air with a heating rate of 10 K per minute. The cylindrical billets (2.06 cm diameter, 0.55 cm thick for S80Z20; 2.54 cm diameter, 0.3 cm  thick for all other compositions) were cut into four quarters for testing. The resulting rounded triangular samples (1.03 cm from the corner to the center of  the arc, 0.55 cm  thick for S80Z20; 1.27 cm from  the corner to the center of the arc, 0.3 cm thick for all other compositions) were rested on quartz glass on the edges of an alumina boat  in order  to minimize  the contact area. All of  the samples were furnace cooled then cross-sectioned and polished for characterization using SEM  (JEOL JSM-6100 and LEO GEMINI 1530). Oxidation  tests were duplicated on separate samples  to verify repeatability of the results.  2.3. Thermal gradient exposure  A schematic of  the  torch  testing set up used  in  this study  is shown in Fig. 2. Before torch testing, the UHTC samples were laid on two MoSi2 rods resting between two refractory bricks. A CH4 -O2 torch (National Torches, Minneapolis, MN) was used to oxidize UHTC samples with temperatures in localized areas capable of  reaching 2373 K. This  temperature was  achieved directly under  the cone of  the  torch;  in addition,  the cone creates a  temperature gradient across  the surface of  the sample.          \\x0c', '3878   Table 2  P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883  Gas-condensed phase equilibria used to calculate ZrB2 -SiC volatility diagram.  Species  B2O3 (g)  B2O2 (g) BO2 (g)  BO(g)  B2O(g)  B2 (g)  B(g)   SiO2 (g)  SiO(g)   Si(g)   Species in equilibrium with ZrB2 -SiC   Species in equilibrium with the stable oxides  ZrB2 + 2.5O2 (g) = ZrO2 + B2O3 (g)  ZrB2 + 2O2 (g) = ZrO2 + B2O2 (g) ZrB2 + 3O2 (g) = ZrO2 + 2BO2 (g) ZrB2 + 2O2 (g) = ZrO2 + BO(g)  ZrB2 + 1.5O2 (g) = ZrO2 + B2O(g)  ZrB2 + O2 (g) = ZrO2 + B2 (g)  ZrB2 + O2 (g) = ZrO2 + 2B(g)  SiC + 1.5O2 (g) = SiO2 (g) + CO(g)  SiC + O2 (g) = SiO(g) + CO(g)  SiC + 0.5O2 (g) = Si(g) + CO(g)   B2O3 (l) = B2O3 (g) B2O3 (l) = B2O2 (g) + 0.5O2 (g) B2O3 (l) + 0.5O2 (g) = BO2 (g) B2O3 (l) = 2BO(g) + 0.5O2 (g) B2O3 (l) = B2O(g) + O2 (g) B2O3 (l) = B2 (g) + 1.5O2 (g) B2O3 (l) = 2B(g) + 1.5O2 (g) SiO2 (l) = SiO2 (g) SiO2 (l) = 0.5O2 (g) + SiO(g) SiO2 (l) = O2 (g) + Si(g)  may reflect  local variations  in  the composition from  the nominal value, but  large-scale deviations were not apparent during SEM/EDS examination. Using energy dispersive X-ray spectroscopy  (EDS),  the dominant phase on  the  surface of each sample after oxidation was verified as SiO2 (Fig. 4) based upon  the Si/O composition  ratio. The SiO2 scale on all compositions was amorphous as judged by the smooth clear appearance without any grains. Each of  the samples experienced minimal mass gain after oxidation  (less  than 0.08 g) which  is expected  for  the passive  Fig. 4. Plan-view BSE micrographs after oxidation at 1773 K for 50 h for (a) S80Z20, (b) S65Z35, (c) S50Z50, (d) S35Z65, and (e) S20Z80.  \\x0c', 'P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883   3879  Fig. 5. Cross-sectional view of (a) S80Z20, (b) S65Z35, (c) S50Z50, (d) S35Z65, and (e) S20Z80 with corresponding Si K␣, O K␣, and Zr L␣  X-ray maps after  oxidation at 1773 K for 50 h.  oxidation of ZrB2-SiC composites. There was no discernible reaction between  the  samples and  the crucibles. There were, however,  important differences  in  the  thicknesses and oxide structures between each of  the compositions. Fig. 5  shows a cross-sectional view of each sample after oxidation as well as the corresponding O K␣, Si K␣, and Zr L␣  X-ray maps. One feature that is immediately apparent is that the thickness of the SiO2 scale decreases with increasing SiC content, except at the highest SiC content. Another important feature in the cross-sectional images can be observed  in  the Si K␣  X-ray maps. In  three of  the samples (i.e. S50Z50, S35Z65, and S20Z80), there is a well defined SiCdepleted layer (Fig. 5(c)-(e)).  (Fig. 6). The depression was most prominent in the S65Z35 sample (Fig. 6(b)) and a smaller depression occurred on the surface of the S80Z20 sample. The three samples on the other end of the ZrB2-SiC pseudobinary  (S50Z50,  S35Z65,  and  S20Z80)  did  not  show  any evidence of a depression on the surface (Fig. 6(b)). Instead, the areas under the hot zone of the torch were void of SiO2 with only a layer of ZrO2 remaining. Fig. 7(c)-(e) shows a cross-sectional view of the three ZrB2 -rich samples as well as the corresponding O K␣, Si K␣, and Zr L␣  X-ray maps. The  regions containing oxides were almost entirely composed of ZrO2 although the S35Z65 and S50Z50 samples contained a small amount of residual SiO2 in the oxide layer.  4.2. Thermal gradient torch study  5. Discussion  The five UHTC compositions were evaluated after a 20 min ≈2073 K under the cone of the exposure to a CH4 -O2 torch at  flame. UHTC samples on both ends of  the ZrB2 -SiC pseudobinary experienced a minimal mass gain (less than 0.1 g) during the extended exposure although SiC-rich  samples  show evidence of depression in the area under the hot zone of the torch  5.1.   Isothermal furnace study  The ZrB2-SiC  volatility  diagram  (Fig.  3),  provides  an effective means  to analyze  the experimental  results. The diagram  indicates  that  the oxidation products of  this composite at 1773-2073 K  should be SiO2 , ZrO2 , and B2O3 . This was  \\x0c', '3880   P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883  Table 3  Reaction layer thicknesses for all compositions oxidized at 1773 K for 50 h.  Sample  S80Z20   S65Z35  S50Z50  S35Z65  S20Z80   SiO2 thickness (\\u242em)   ±  ± ±  490   50  65 ± 100 ±  10   10  120   10   120    10   SiC-depletion (\\u242em)  None  None  80    20  ± ± ±  170   260    10   75  depletion  layer  increased by 20  \\u242em and 180  \\u242em, respectively, when the ZrB2 content was increased from 50 to 80 vol%. The widening SiC-depleted  layer with  increasing ZrB2 content can be  justified by  the available SiC reservoir under  the protective SiO2 scale. Not only did  the S20Z80 sample has  the  thickest SiO2 scale, but  it also has  the  least amount of SiC  to support the growth of SiO2 (l). The samples rich in SiC (i.e. S65Z35 and S80Z20) have a comparatively large SiC reservoir and a thinner SiO2 scale  to support. As a  result, very  little SiO(g)  is  transported  from  the SiC-depleted  layer  to  the SiO2 surface  layer which explains the absence of a discernible SiC-depleted layer in  the X-ray map of  the S65Z35 sample. Again,  this reasoning supports the conclusions made by Fahrenholtz.4 The  thickness  trend of  the SiO2 layer  is compared  to  the sample composition in Table 3. Over the range of compositions examined the trend indicates a general decrease in oxide thickness with increasing SiC content except for the S80Z20 sample. This trend is similar to that observed by Chamberlain et al.24 and Karlsdottir and Halloran34 for ZrB2 -rich compositions. During the initial transient period, as full coverage by the borosilica surface layer develops, there are two competing factors influencing the oxidation  response. For ZrB2 -rich compositions  there  is a reduced amount of Si  to develop SiO2 during oxidation, but the  resulting silica has a  lower viscosity due  to a high B2O3 content as shown by Karlsdottir and Halloran.34 The opposite conditions hold for the SiC-rich compositions. The experimental observations clearly  indicate  that  the composite SiC composition is the main factor controlling the initial oxidation behavior. Moreover, beyond  the  initial period  it  is  important  to note  that based upon  the Fahrenholtz model4 the SiO2 (l)  layer grows at the SiO2 /ZrB2 + SiC-depletion interface as a result of inward diffusion of oxygen and outward diffusion of SiO(g). The greater thickness of  the SiO2 (l)  layer  in ZrB2 -rich composites  is  the result of a greater volume of substrate being captured in the  gradient that causes active oxidation of SiC to SiO(g) + CO(g). This condition is satisfied experimentally by the increasing SiCdepleted layer thickness with increasing ZrB2 content (Table 3 and Fig. 5). Thus, a greater volume of SiO(g)  is  transported across the depleted layer and oxidizes in the SiO2 (l) layer causing it to grow. The experiments conducted at 1773 K also suggest the importance of  the  initial boron content  in  the composite. While  it  is understood  that  the majority of  the B2O3 will evaporate at  this temperature,5 the  large supply of boron at  the glass/substrate interface for ZrB2 -rich UHTCs provides a significant  increase in  the fluidity of  the continuous glass as demonstrated by  the observation of  convection. Additionally,  the Stokes-Einstein  pO2  Fig. 6. Optical view after   torch   testing   for 20 min with   (a) S80Z20 showing  sample recession in the hot zone and (b) S20Z80 showing no depression in the  hot zone.  verified experimentally for SiO2 and ZrO2 at 1773 K (Fig. 4) for all ZrB2-SiC compositions. Because B  is a  light element,  it  is not possible to detect B2O3 reliably by EDS. Moreover, zircon (ZrSiO4 ) is known to be stable in the temperature range of interest, but this phase was not found in the oxidation products in the current work or in other reported work at room temperature and has not been included in Fig. 3.4,33 The existence of a SiC-depleted layer indicates that compo≤50 vol% will form a SiC-depleted sitions with SiC content of  layer during an  isothermal oxidation  test at 1773 K. According  to  the  thermodynamic  analysis  done  by Fahrenholtz,  a SiC-depletion should occur in all ZrB2-SiC compositions,4 but depends on  the  level at  the  interface. For  the S65Z35 and S80Z20 samples,  the  lack of a discernable SiC-depletion  layer can be related  to  the  low mobility of oxygen  in  the continuous SiO2 layer. Assuming  the oxidation mechanism suggested by Fahrenholtz,4 where SiO2 grows by the transport of SiO(g), the amount of silicon (Si) in the SiO2 surface layer is derived from the depletion  layer. Table 3 shows  the  thickness of  the surface oxide layer and corresponding SiC-depleted layer for each of the five samples along the ZrB2 -SiC pseudo-binary. Both the thickness of the SiO2 scale and depletion layer gradually increased as the ZrB2 content increased. On average, the SiO2 scale and the  pO2  \\x0c', 'P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883   3881  Fig. 7. Cross-sectional view of (a) S80Z20, (b) S65Z35, (c) S50Z50, (d) S35Z65, and (e) S20Z80 with corresponding Si K␣, O K␣, and Zr L␣  X-ray maps after ≈2073 K for 20 min.  torch testing at   relation can be used  to show  that  the diffusion coefficient of oxygen through borosilica glass is significantly larger than that though pure SiO2 (≈10 −12 m2 /s and  ≈10 −21 m2 /s for pure B2O3 and pure SiO2 , respectively).7 This will result in a thicker glass which inevitably requires a larger supply of Si to the surface.  5.2. Thermal gradient study  The existence of a depression  in  the  (S65Z35 and S80Z20) can be attributed  SiO2 by the reaction:  two SiC-rich samples to  the evaporation of  )  x  p  (  g o  l   SiO(g)    1/2O2 (g)   (2)  ⇒  SiO2 (l)   +  This reaction occurs in the hot zone of the torch. The gradient tests were conducted such  that  the hot zone of  the sample was about 2073 K while  the outer edges of  the sample were about 1873 K. The volatility diagram for the oxidation of SiC (Fig. 8), shows that SiO(g) becomes the predominant vapor phase in equilibrium with  the condensed oxide above about 2073 K (see  the in air” line in Fig. 3). There will be enhanced evaporation of the SiO2 protective layer above the transition. Therefore, it is  ‘‘pO2  10  5  0  −5  −10  −15  −20  SiC  SiO2()l  SiO(g)  1500°C 1650°C 1800°C  −30  −20  −10  log(pO2  )  SiO2(g)  r  i  a  n  i  2  O  p  0  10  Fig. 8. SiC volatility diagram at 1500     C, 1650     C, and 2073 K.    \\x0c', '3882   Table 4  P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883  the predominant vapor phase  in equilibrium with SiO2 (l) near 2073 K. This will cause preferential evaporation of SiO2 (l)  to SiO(g). For ZrB2-SiC composites oxidized at 2073 K, ZrO2 is the only stable oxide; therefore, samples rich in ZrB2 will grow a ZrO2 skeleton initially. With continued exposure the ZrO2 may form a compact  layer, but even  fully densified ZrO2 demonstrates a high parabolic rate constant at 2073 K.12 At  the same time, a recent computational study by Li et al.25 suggested that a relatively small amount of SiC (i.e. 20 vol% SiC), as opposed to pure ZrB2 , may be advantageous at  temperatures above  the SiO2 (l) to SiO(g) transition temperature. The oxidation performance was modeled by a “solid pillar, liquid roof” construction that may help resist spallation of the oxide. Based upon the hypothetical construction, it was claimed that the ZrO2 “pillars” can actually help avoid shearing of the SiO2 (l) layer due to capillary forces between the pillars. However, at ultra-high temperatures, the rapid growth kinetics of ZrO2 does not provide sustainable long  term oxidation protection.12,20 Indeed,  there needs  to be a concurrent future effort in mitigating the diffusion of oxygen through  the glass as well as  through  the ZrO2 for ultra-high temperatures. Nevertheless, below  the SiO2 (l)  to SiO(g)  transition  temperature,  there  is a clear advantage  in  reducing  the ZrB2 content below 80 vol% so as  to retain a continuous SiO2 matrix in the mixed scale and to enhance the oxidation resistance considerably.  6. Conclusions  All ZrB2 -SiC composites from 20 to 80 vol% SiC form a continuous layer of SiO2 after oxidation in air for 50 h at 1773 K. For compositions from 20 to 80 vol% SiC there is a systematic trend for  the SiO2 layer  thickness  to decrease. Over  the same composition range the SiC-depletion layer thickness also decreases systematically and becomes undetectable at 65 vol% SiC. A thermal gradient torch study was conducted at a temperature of about 2073 K under the cone of the torch with the edges of the UHTC samples reaching about 1873 K. SiC-rich compositions (i.e. S80Z20 and S65Z35) experienced sample depression in  the hot zone of  the  torch. This behavior was analyzed using the SiC volatility diagram which indicates that SiO(g) becomes the predominant vapor phase above about 2073 K as opposed to SiO2 (g)  lower  temperatures. At 2073 K, ZrB2 -rich compositions  (i.e. S35Z65 and S20Z80) did not  form a depression, rather a ZrO2 skeleton developed where SiO2 (l) was lost due to SiO(g) evaporation. At ultra-high temperatures sample processing irregularities can influence the initial oxidation response, but this  is of  little consequence  to  the  longer  term response where the loss of SiO2 and the residual, ZrO2 control the performance. Finally,  the  results  from  this study have demonstrated  that the UHTCs  investigated  in  the current available  literature (i.e. 10-30 vol% SiC) are not necessarily  the optimal composition for high  temperature  (<1773 K) performance. For  isothermal exposure it was found that the SiC-depleted layer can be reduced by increasing the volume fraction of SiC, and is undetectable in composites with at  least 65 vol% SiC. Also,  the overall weight density of the ceramic can be reduced with improved oxidation performance at 1773 K by  increasing  the SiC content. Under  SiC-depleted   layer   thickness for all compositions   torch   tested at   ≈2073 K for  20 min.  Sample   S80Z20  S65Z35  S50Z50  S35Z65   S20Z80   SiC-depletion (\\u242em)  None  None  ± ± ±  110   160   170    10   15   10  likely  that  the SiO2 liquid  in  the hot zone of  the  torch reaches the SiO(g) transition and evaporates causing a depression on the surface of  the sample. It also must be noted  that  there  is water vapor present during torch testing:  CH4 (g)   +   2O2 (g)   ⇒   CO2 (g)   +   2H2O(g)   (3)  Nguyen  et  al.23 showed  that  the  addition of water vapor during ultra-high  temperature oxidation can cause additional volatile species in ZrB2-SiC composites to form (e.g. Si(OH)4 ) as compared  to  the volatile species expected during static air furnace testing. The  three samples on  the other side of  the binary (S50Z50, S35Z65, and S20Z80) displaying a small amount or absence of SiO2 and a noticeable  layer of ZrO2 suggests  that  the volume fraction of ZrB2 in these samples was large enough to support a ZrO2 skeleton that grows in the region of evaporating SiO(g). Another intriguing feature found in the X-ray maps (Fig. 7) is  the development of  the SiC-depleted  layer. Like  the  results from the isothermal oxidation tests at 1773 K, only compositions that contained at least 50 vol% ZrB2 formed a discernable SiCdepleted  layer. The S35Z65 and S50Z50 samples have clearly defined regions of SiC-depletion and  the depleted  layer  in  the S20Z80 sample can be found in the region that contains no oxides and no Si. This behavior occurred  in  the S20Z80 composition because of the small volume fraction of SiC in the original sample. Pure SiO2 will evaporate above about 2073 K leaving only a ZrO2 oxide layer and a SiC-depleted layer in the S20Z80 sample. Samples with greater than 20 vol% SiC retained some SiO2 in  the oxide  layer after 20 min; however,  it  is expected  that all the SiO2 will evaporate in all compositions after longer exposure times. Table 4 shows the thickness of the SiC-depleted layer in each of the compositions. Oxide thicknesses were difficult to quantify for  torch  tested samples due  to  the steep  thermal gradient and substrate recession during the ultra-high temperature exposure. The depletion layer, however, was found to consistently increase with increasing ZrB2 content which is identical to the behavior seen  in  the  isothermal  furnace  tests. The zirconia  formation, however,  is  thicker  in  the SiC-rich composite presumably due to the longer transient stage required to form a continuous silica layer over the zirconia structure. ≈2073 K, the limited surIn this case, where temperatures are  face depression appears to indicate some advantage of a reduced 104 Pa SiC content. At 2073 K the vapor pressure of B2O3 is  but the vapor pressure of ZrO2 is still very low (10 −5 Pa). Also, the SiC volatility diagram (Fig. 8) shows  that SiO(g) becomes  \\x0c', 'P.A. Williams et al. / Journal of the European Ceramic Society 32 (2012) 3875-3883   3883  rapid heating exposure in a torch, the high SiC content (S80Z20) composite exhibited depression in the hot zone, but the integrity of the underlying microstructure was maintained for the duration of the 20 min test.  Acknowledgements  The  support  from  the National Aeronautics  and  Space Administration  (Cooperative Agreement: NNX08AB35A)  for the  initial stage of research, and  the continuing support of  the Air Force Office of Scientific Research (Grant number: FA955011-1-0201-AFOSR) is gratefully acknowledged.  References  1. Wuchina E, Opila E, Opeka M, Fahrenholtz W, Talmy   I. UHTCs: ultra high   temperature ceramic materials for extreme environment applications.  Electrochem Soc Interface 2007;16:30-6.  2. Fahrenholtz WG, Hilmas GE. Oxidation of ultra-high temperature transition metal diboride ceramics. Int Mater Rev 2012;57:61-72.  3. Dimiduk DM, Perepezko JH. Mo-Si-B alloys: developing a revolutionary turbine engine material. MRS Bull 2003;28:639-45.  4. Fahrenholtz WG. Thermodynamic analysis of ZrB2 -SiC oxidation: formation of a SiC-depleted region. J Am Ceram Soc 2007;90:143-8.  5. Karlsdottir SN, Halloran JW, Henderson CE. Convection patterns in liquid  oxide films on ZrB2 -SiC composites oxidized at a high temperature. J Am Ceram Soc 2007;90:2863-7.  6. Karlsdottir SN, Halloran JW. Rapid oxidation characterization of ultra-high temperature ceramics. J Am Ceram Soc 2007;90:3233-8.  7. Karlsdottir SN, Halloran   JW. Formation of oxide   scales on   zirconium  diboride-silicon carbide composites during oxidation: relation of subscale recession to liquid oxide flow. J Am Ceram Soc 2008;91:3652-8.  8. Han   J, Hu P, Zhang X, Meng S, Han W. Oxidation-resistant ZrB2 -SiC composites at 2200     C. Compos Sci Technol 2008;68:799-806.  9. Eakins   E,   Jayaseelan  DD,  Lee WE.  Toward  oxidation-resistant temperature ceramics. Met Mater Trans 2011;42A:  ZrB2 -SiC ultra high  878-87.  10. Zhang XH, Hu P, Han  JC. Structure evolution of ZrB2 -SiC during  oxidation in air. J Mater Res 2008;23:1961-72.  the  11. Cutler RA. Engineering properties of borides. In: Schneider Jr SJ, editor.  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Oxidation-based materials selection for  2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci 2004;39:5887-904.     19. Opeka MM, Talmy   IG, Wuchina EJ, Zaykoski JA, Causey SJ. Mechani cal, thermal, and oxidation properties of refractory hafnium and zirconium compounds. J Eur Ceram Soc 1999;19:2405-14.  20. Monteverde F, Bellosi A. Oxidation of ZrB2 -based ceramics  Electrochem Soc 2003;150:B552-9. 21. Ordan `yan SS, Dmitriev AI, Moroshkina ES. Reaction of SiC with ZrB2 . Inorg Mater 1989;25:1487-9.  in dry air. J  22. Clougherty EV, Pober RL, Kaufmann L. Synthesis of oxidation  resistant metal diboride composites. Trans Metall Soc AIME 1968;242:1077-82.  23. Nguyen QN, Opila EJ, Robinson RC. Oxidation of ultrahigh  temperature ceramics in water vapor. J Electrochem Soc 2004;151:B558-62.  24. Chamberlain A,   Fahrenholtz W, Hilmas G, Ellerby D. Oxidation   of  ZrB2 -SiC ceramics under atmospheric  Trans 2005;1:1-8.  reentry conditions. 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Texas: Chemistry Software   [cited 05.03.2010;  last accessed 29.11.2011]. Available from: http://www.hsc-chemistry.net/.  32. Heuer AH, Lou VLK. Volatility diagrams   for   silica,   silicon-nitride, and  silicon-carbide and their application to high-temperature decomposition and oxidation. J Am Ceram Soc 1990;73:2789-803.  33. Lou VLK, Mitchell TE, Heuer AH. Review-graphical   displays   of   the  thermodynamics of high-temperature gas-solid   reactions and   their appli cation  to oxidation of metals and evaporation of oxides. J Am Ceram Soc 1985;68:49-58.  34. Karlsdottir SN, Halloran   JW. Oxidation of ZrB2 -SiC:  liquid  oxide  phase  formation.   influence of SiC  content  on  solid  2009;92:481-6.  and   J Am Ceram   Soc  \\x0c']"
},{
  "_id": 182,
  "PDF": "Oxidation of ZrB2-SiC- Influence of SiC Content on Solid and Liquid Oxide Phase Formation.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 92 [2] 481 - 486 (2009)  DOI: 10.1111/j.1551-2916.2008.02874.x  r 2009 The American Ceramic Society  Oxidation of ZrB2-SiC: Influence of SiC Content on Solid and Liquid  Oxide Phase Formation  Sigrun N. Karlsdottir  w,z  and John W. Halloran  Department of Materials Science and Engineering, University of Michigan, Ann Arbor, Michigan 48104  phases  The effect of SiC concentration on the liquid and solid oxide of ZrB2-SiC composites features called convection cells are  investigated. Oxide-scale  oxidation  formed  during  is  idation  formed from liquid and solid oxide reaction products upon oxthe ZrB2-SiC composites. These form in the outermost borosilicate oxide film of the oxide scale  convection  cells  of  formed on the ZrB2-SiC during oxidation at high temperatures (\\x15 15001C). In this study, three ZrB2-SiC composites with different amounts of SiC were tested at 15501C for various durations of time to study the effect of the SiC concentration par ticularly on the  formation of  the  convection cell  features. A  calculated ternary phase diagram of a ZrO2-SiO2-B2O3 (BSZ) system was used for interpretation of the results. The convection  cells  formed during oxidation were  fewer and less uniformly  distributed for composites with a higher SiC concentration. This is because the convection cells are formed from ZrO2 precipitates from a BSZ oxide liquid that forms upon oxidation of the composite at 15501C. Higher SiC-containing composites will have less dissolved ZrO2 because they have less B2O3, which results in a smaller amount of precipitated ZrO2 and consequently fewer convection cells.  I.  Introduction  THE boride-based composites, ZrB2-SiC and HfB2-SiC, are considered among the most promising ceramic composites for high temperature and structural applications due to their unique properties.1 The ZrB2-SiC and HfB2-SiC composites are ultra-high-temperature ceramics, that are oxidation resistant at  high temperatures due to the presence of a complex multilayer oxide scale that slows down inward diffusion of the oxygen.1-4 It the oxidation behavior4 by  is well known that SiC improves  adding silica to the oxide film, but the details of the role of SiC in  the amounts and compositions of the phases in the film have not  been described in detail. This paper compares the oxide film microstructure after oxidation at 15501C for three different SiC  concentrations. During high-temperature oxidation (412001C) of the ZrB2- SiC composite, a liquid oxide film, borosilicate (SiO2-B2O3), forms on the outer surface.5,6 Solid zirconium oxide (ZrO2) is also formed along with the liquid oxide film above the unreacted  ZrB2-SiC material. Because of the high vapor pressure of boria (B2O3(l )) at these temperatures, compared with silica (SiO2(l )), the B2O3 is preferentially evaporated from the borosilicate liquid. The liquid oxide film at the outer surface then becomes a predominantly viscous SiO2-rich borosilicate liquid.1,2,7-9 A novel oxidation mechanism involving flow of the proposed.10 Flow of  boria-silica-zirconia  liquid  oxide  been  has  a  (ZrO2-SiO2-B2O3 (BSZ)) liquid was used to explain the distinc M. Cinibulk—contributing editor  Manuscript No. 24488. Received March 31, 2008; approved November 5, 2008. zCurrent address is at the Department of Materials, Biotechnology and Energy, Innovation Center Iceland, IS-112 Reykjavik, Iceland.  Author to whom correspondence should be addressed. e-mail: nanna@umich.edu  w  tive microstructural features on the external oxide surface and in cross section.11 These features, called convection cells, consist of  ZrO2 ‘‘islands’’ located in larger SiO2-rich ‘‘lagoons’’ with B2O3rich patterns surrounding the islands.10-12 These features are  shown in Fig. 1, which shows a surface of a ZrB2-15 vol% SiC composite oxidized at 16001C for 30 min. The area around the  convection cells consists of a SiO2-rich glass with small micrometer-sized zirconia dispersoids. The B2O3-rich flower petal-like patterns are visible in backscattered electron (BSE) imaging, and  in stronger contrast in cathodoluminesence imaging, but are not  observed in secondary electron imaging scanning electron mi croscopy (SEM). We previously suggested that ZrO2 cores form by precipitation during the evaporation of boria (B2O3) from a BSZ liquid that rises through the outer SiO2-rich borosilicate layer and flows laterally, forming the B2O3-rich regions (the petals) around the ZrO2 islands.10,11 The BSZ liquid is formed when a borosilicate liquid rich in B2O3 dissolves some ZrO2. We estimate the compositions of these phases with the aid of a  calculated isothermal section of a B2O3-SiO2-ZrO2 system at 15001C to describe the equilibrium between a BSZ liquid and crystalline ZrO2.11 A distinction is made between two zirconia morphologies formed during oxidation. The ‘‘primary’’ zirconia  (porous underlying ZrO2(s)) is formed directly by the oxidation of ZrB2 by oxygen diffusion through the SiO2-rich borosilicate primary surface layer. The ‘‘secondary’’ zirconia is precipitated  from the BSZ liquid. At the surface, the B2O3 evaporates from the BSZ liquid and secondary zirconia precipitates. The surface  is  then covered with the less volatile phases  i.e.  the silica-rich  liquid and secondary zirconia, which are located near the site of  evaporation (near  the B2O3 the B2O3 petals). The zirconia precipitates form the zirconia island and dispersed zirconia.11 What we observe at room temperature are patterns of  secondary  crystalline zirconia and glasses cools.10-12  formed after  the silicate liquid  The convection cell theory suggests that the formation of the  convection cells  is dependent on the composition of  the BSZ  liquid, which in turn is dependent on the composition of  the  ZrB2-SiC composite. Here,  three ZrB2-SiC composites with  Fig. 1. Backscattering electron image of a surfaces of a ZrB2-15 vol% SiC composite tested at 16001C for 1 2 h, showing an example of convec tion cells and their patterns.  481  \\x0c', 'different SiC concentrations are tested at 15501C for various  durations of  times and compared with study the effect of SiC  concentration on the formation and distribution of convection  cells.  II.  Experimental Procedure  Three ZrB2-SiC-based tions, 15, 20, and 30-vol% SiC, were used for  composites with  different  concentra this study. We  presume that  the decisive difference between the composites is  their SiC concentration. These were provided by different  lab oratories and differed slightly in the processing, as described  below. The ZrB2-15 vol% SiC (ZS15) composite used for study was provided and fabricated by Dr. Alida Bellosi and  the  Dr. Frederic Monteverde at The Institute of Science and Tech nology for Ceramics, National Research Council (ISTEC-CNR)  in Faenza, Italy. The properties and the processing of the ZS15 composite are presented in more detail elsewhere.13 The ZrB2 composites containing the 20% SiC and 30 vol% SiC were  fabricated and provided by Dr. William G. Fahrenholtz, Dr.  Gregory E. Hilmas, and associates at  the University of Mis souri-Rolla (UMR) in Rolla, MI. The ZrB2-20 vol% SiC (ZS20) composite contained around 2 vol% WC that was incorporated  in the attrition milling step during fabrication of the material at  UMR due to the use of WC milling media. The fabrication and materials are described in more detail elsewhere.14 The ZrB2-30 vol% SiC (ZS30) composite was attrition milled at UMR with  ZrO2 milling media during fabrication; other than SiC and ZrB2 was not reported, tamination from the milling media. The fabrication and starting materials are described in detail elsewhere.15 Figure 2 shows the  the presence of phases  indicating no con microstructure of the three composites ZS15, ZS20, and ZS30.  Oxidation was conducted in a high-temperature box furnace (SentroTech Corporation, Berea, OH) in ambient air at 15501C  at different dwelling times, ranging from 3 to 8 h. The heating and cooling rates used were 131C/min. The three composites (ZS15, ZS20, and ZS30) were tested at 15501C for 3, 4, and 8 h.  The specimens were supported by sacrificial  support pieces of  the same ZS material. The sacrificial supports were placed on an  Al2O3 support in an Al2O3 crucible. The ZS15 composite was cut into thin sheets from a bulk material with a wire electrical  discharge machine (w-EDM) (Ann Arbor Machine Model 1S15,  Ann Arbor, MI). The thin sheets of the ZS15 were then cut with  a diamond saw (IsoMet  s  1000 diamond precision saw, Buehler,  Lake Bluff, IL) into small rectangular coupons with a total surface area between 0.5 and 1 cm2. Then, ca. 200 mm was removed  from the surface of the ZS15 coupons with a fine diamond grid  (Omni Brade, TBW Industries, Furlong, PA)  to remove any  heat-affected zone that could have formed on the surface of the  ZS15 sheets after the w-EDM. The ZS20 and ZS30 composites  obtained from UMR as tests bars where ground and cut with  the same tools as the ZS15 to obtain the same surface finish and surface area (B0.5-1 cm2) as  a similar  the  specimens  tested.  Before testing, the specimens were ultrasonically acetone and dried at 1001C.  cleaned in  Microstructural analyses were performed on the surfaces and  cross sections using SEM and BSE microscopy. The cross sec tions of  the oxidized specimens were prepared for microstruc tural analysis by nonaqueous polishing procedures down to a 1 mm finish. Specimens were coated with carbon before the  microstructural analysis.  III.  Results: Microstructure Analysis  The purpose of testing other ZrB2-SiC composites with different SiC concentrations was twofold: (1) to verify that convection  cells form in ZrB2-SiC materials other than the ZrB2-15 vol% SiC (ZS15) composite previously studied by the authors10-12,16  and (2) to study the effect of SiC concentration on the formation  and distribution of convection cells. The evaluation of convec tion cells with time and temperature, formed at the oxide scale  during high-temperature oxidation on the ZS15 composite, has been studied in detail and reported elsewhere.16  (1)  Surface Analysis  Microstructure analysis with SEM and BSE imaging revealed that the ZS15, ZS20, and ZS30 specimens tested at 15501C for 3,  4, and 8 h all showed the formation of convection cells on the  surfaces. The convection cells on the ZS20 and ZS30 specimens  were comparable to ZS15. However,  the distribution and the  number of cells formed on the surfaces of  the three specimens  were somewhat different from each other. Firstly, the ZS20 and  ZS30 specimens had,  in general, fewer convection cells than the  ZS15 and they were not as uniform as the convection cells for the ZS30 specimen tested at 15501C for  the ZS15. Secondly,  shorter dwelling times  (3 and 4 h) had smaller ZrO2 sometimes surrounded by B2O3 circles instead of distinct petals. An example of this is shown in Fig. 3, where BSE images are  islands,  shown of the surfaces of the three composites, ZS15, ZS20, and tested at 15501C for 4 h. The images were taken at ZS30, low magnifications (\\x02 36-38, 20 KV, spot size: 5) to be able to compare the population and the distribution of the convection cells  on the three composites. The micrographs in Fig. 3 indicate that  the lower SiC-containing composite (ZS15) has a larger number  of convection cells spread over the surface than the higher SiC containing composites (ZS20 and ZS30). To investigate this, the  number of convection cells was quantified for the ZS15, ZS20,  and ZS30 specimens by visual counting. The number of convection cells per square mm (cell density: no./mm2) was quan tified by counting the number of cells in a blinded fashion from 15 random fields (a square frame; 1 mm \\x02 1 mm) per sample at \\x02 36-38 microscopic magnification. The cell counts were then averaged for each sample. The quantification technique was  based on protocols for manual quantification of biological features.17,18 Figure 4 shows a graph of the average cell densities and the standard error (7standard deviation) for the three composites tested at 15501C for 3, 4, and 8 h. The standard de viations  (the error bars) of  the values for  the ZS20 and ZS30  composites  are  rather  large,  and  the  values  for  the  8-h oxidation time for the ZS20 and ZS30 do not differ much  from ZS15. Overall,  the  results  indicate  that with increasing  amount of SiC, fewer convection cells form. The large standard  deviation for  the ZS20 and ZS30 composites  (in Fig. 4) was  most likely because the convection cells on the ZS20 and ZS30  Fig. 2.  Secondary electron images of the microstructure of the three different composites tested (a) ZrB2-15 vol% SiC (ZS15), (b) ZrB2-20 vol% SiC (ZS20), and (c) ZrB2-30 vol% SiC (ZS30). The dark phases are the SiC grains and the light gray are the ZrB2 grains, additionally the white phase for the ZS20 is reported elsewhere to be WC grains.14  482  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 92, No. 2  \\x0c', 'February 2009  Oxidation of ZrB2-SiC: Influence of SiC Content on Oxide Phase Formation  483  Fig. 4.  The density of convection cells  (no.  cells/mm2)  for  the  three  composites containing different amount of SiC; ZS15, ZS20, and ZS30 tested at 15501C for 3, 4, and 8 h.  convection cells formed on the three composites is also indicated  by ZrO2 dendrites that were observed close to the ZrO2 islands for all three composites. The formation of these dendrites is discussed in detail elsewhere.16 An example of  these dendrites  on the ZS15, ZS20, and ZS30 composites  is  shown in Fig. 5,  where BSE images of the surfaces of ZS15, ZS20, and ZS30, tested at 15501C for 3, 3, and 8 h, respectively, are shown. The  presence of dendritic zirconia adjacent to the islands is presented  as evidence that  the secondary zirconia is precipitating from a  saturated liquid solution.  (2)  Cross-Sectional Analysis  Figure 6 shows  the  secondary electron micrographs of  cross  sections of the ZS15, ZS20, and ZS30 specimens tested for 8 h. at 15501C. The micrographs  show enhanced oxidation zones  (deeper recession) under the ZrO2 islands where the cross sections cut through the convection cells. The contrast conditions  for Fig. 6 do not clearly distinguish the interface between the  oxide and the unreacted bulk, but  these are shown clearly in  backscattered imaging, and we  indicate  the  interface using a  broken line. The enhanced oxidation zones under the ZrO2 islands of the convection cells are discussed elsewhere.12 But, in  summary,  it forms when the BSZ liquid squeezes up to the sur face and thus a path is open for more rapid inward diffusion of  oxygen due to the lower viscosity of the BSZ liquid compared with the surrounding SiO2-rich surface liquid layer.12 The thickness of the oxide layer for each composite increases  with increasing dwelling time;  this  is  shown in Fig. 7, which  shows graphs of the total thickness of the oxide layer versus dwelling time at 15501C for the three composites. The total ox Fig. 3.  Backscattering electron images of the surfaces of the ZrB2 composites containing different amount of SiC; ZrB2-15 vol% SiC (ZS15), ZrB2-20 vol% SiC (ZS20), and (c) ZrB2-30 vol% SiC (ZS30), all tested at 15501C 4 h.  specimens were not uniformly distributed; thus, the cell density  in one area differed from other areas. The decrease in the cell  ide thickness was calculated as the sum of the measured averages  density in the ZS15 composite for the longer oxidation time at 15501C,  i.e. 8 h, can be seen in Fig. 4. It was previously shown  that  some  convection cells may become  ‘‘extinct’’ after  they  form and later become covered by fresh SiO2 liquids depositing on the surface. Thus, the cell density decreases with longer oxidation time.16  of the SiO2-rich borosilicate surface layer and the underlying ZrO2 layer ðdavg ¼ dSiO2 rich þ dZrO2 Þ. The thickness of each layer was measured by averaging 20-30 values at various locations. A  large standard deviation in the oxide thickness was found due to  the deeper recession under the convection cells.  Fewer convection cells were found in the cross sections of the  In general, the features of the convection cells formed on the  higher SiC-containing composites. This  is  in agreement with  ZS20  and ZS30  are  very  similar  to what was observed on  the surface analysis; there was fewer convection cells as the SiC  the ZS15 specimens,  i.e. ZrO2 islands surrounded by B2O3-rich (petals) and outer SiO2 lagoons. The similarity of the  regions  content increased. From Figs. 6 and 7 it is evident that there is  less oxidation for the higher SiC-containing composites. This is  Fig. 5.  Backscattering electron micrographs of a ZrO2 islands on the surfaces (a) ZS15, (b) ZS20, and (c) ZS30 composites tested at 15501C for 3 h (a), (b) and 8 h. (c), respectively, showing dendrites (indicated by arrows) grown close to the ZrO2 islands.  \\x0c', 'consistent with what has been published by Chamberlain et al.14  on the ZrB2-20 vol% SiC composite material and other ZrB2- SiC composites with various SiC concentrations tested under  similar oxidation conditions.  IV.  Discussion and Phase Equilibrium Analysis  The results of the present study show that convection cells form,  in general on ZrB2-SiC composites during oxidation at higher (B15501C). This supports temperatures the previous hypothesis10,11 that the formation of the convection cells is dependent on  the formation of a transient BSZ liquid and a large volume in crease that occurs during oxidation of ZrB2-SiC-based materials. Fewer convection cells (no. of cells/mm2) form on a higher  SiC-containing composite,  i.e. ZS20 and ZS30, compared with  ZS15, the lower SiC-containing composite. We suggest that this  may be due to the difference in the composition of the BSZ liq uid that  forms during oxidation. The ZS20 and especially the  ZS30 show less B2O3 (in the BSZ(l )) formation compared with the ZS15; thus, there should be less ZrO2 dissolved in the BSZ liquid, which should result in formation of fewer ZrO2 islands. This is supported by a calculated ternary phase diagram for the B2O3-SiO2-ZrO2 system at 15001C published previously by the authors and Karlsdottir et al.,11 and from the microstructural  analysis and quantification of convection cells shown here.  The calculated ternary phase diagram for the B2O3-SiO2- ZrO2 system at 15501C does not exist and so we have to use the only available ternary phase diagram for this system, which is a calculated phased diagram at 15001C (the experimental diagram  does not exist either). We justify this by pointing out that from a calculated binary phase diagram of the B2O3-ZrO2 system,11 the solubility of the zirconia would only be a slightly higher at 15501C. Thus, we are not overestimating the amount of second ary zirconia. The authors do acknowledge that  the calculated  ternary phase diagram must be regarded as somewhat specula tive. However, the general features of the system can be expected  to be qualitatively correct. A detailed description of the calcuthe phase diagram is given elsewhere.11 It should be  lation of  stated here  that  although the  calculated phase diagram can  be used for  the interpretation of  the results,  there is a strong  need for  experimental phase diagram for  the  ternary B2O3- SiO2-ZrO2 system, which, hopefully, future experimental investigations may provide.  Figure 8 shows the isothermal section of the B2O3-SiO2-ZrO2 system at 15001C, annotated on the lower right to show the  equilibrium compositions  expected for  the  fresh products of  oxidation of zirconium diboride with 15, 20, or 30 vol% SiC.  These oxide compositions fall  in the two-phase region with the  BSZ liquid solution and crystalline zirconia (ZrO2). The phase diagram in Fig. 8 shows (indicated by arrows) how higher SiC containing material would have a smaller amount of dissolved 15001C.  ZrO2 This is because the B2O3 dissolves the ZrO2 and SiO2 but higher SiC-containing composites, there is relatively less B2O3 and thus there will be less dissolved ZrO2 in the BSZ liquid.11 Annotations on the upper left portion of Fig. 8 also show the  in the BSZ liquid formed upon oxidation at  for  compositions expected after the B2O3 evaporates from the BSZ liquids (assuming 90% of the B2O3 has evaporated away as an example).  Consider the oxidation of a ZrB2-15 vol% SiC composite at 15001C. We expect the phase assemblage for the complete ox idation of the composite to be, on a molar basis, 0.33 mol solid  ZrO2 and 0.67 mol of a BSZ liquid (liquid composition 71 mole% B2O3118 mol% SiO2111 mol% ZrO2; this composition is indicated on the phase diagram with a white circle, see Fig. 8).  At these temperatures, the B2O3(l ) will preferentially evaporate when the BSZ liquid reaches the surface and the dissolved ZrO2 must precipitate out. If, for example, 90% of the B2O3(l ) evaporates away, assuming it stops due to increased viscosity when  the relative amount of SiO2 increases in the BSZ liquid and only 10% of B2O3 remains in the liquid, then there would be a shift towards the two-phase region in the phase diagram where ZrO2 precipitates, and the final equilibrium phase assemblage would  be a two-phase mixture, solid ZrO2 (precipitated) and SiO2-rich BSZ liquid, with the BSZ liquid composition being 88 mol%  SiO2, 2 mol% ZrO2 and 10 mol% B2O3, and with the composition changing to 58 mol% SiO2, 36 mol% ZrO2, and 6 mol% B2O3. The equilibrium phase assemblage at this composition is 66 mol% of a SiO2-rich liquid and 34 mol% solid precipitated ZrO2. The amount of ZrO2 that will precipitate under these oxidation conditions for the higher SiC concen total  trated composites (ZS20 and ZS30) can be estimated in the same  way as in the phase diagram, and is shown in Table I.  In Table I,  the estimated amount of precipitated ZrO2(s) the three different ZrB2-SiC composites. It also shows how the solubility of the ZrO2(s) in the BSZ liquid decreases with increasing viscosity and SiC concentration (i.e.,  is  shown for  in creasing SiO2 amount). The viscosity of was estimated from a relationship extrapolated from data by Jabra et al.19 discussed in detail elsewhere.10,11 The estimated  the borosilicate melt  amount of precipitated ZrO2 decreases with increasing SiC content in the composites as shown in Table I. For example the  ZrB2-15 vol% SiC composite (ZS15) is estimated to precipitate 22 mol% more ZrO2(s) than the ZrB2-30 vol% SiC (ZS30), when 90 mol% of the B2O3 has evaporated from the BSZ liquid. Less ZrO2, precipitated from the BZS liquid, will result in fewer convection cells formed during oxidation; this could explain  the decrease in the number of convection cell on the ZrB2-SiC  Fig. 6.  Secondary electron micrographs of the cross-sections of (a) ZS15, (b) ZS20, and (c) ZS30 composites tested at 15501C for 8 h. The interface  between the ZrO2 layer and the unreacted bulk material  is indicated with a broken line. The darker outermost layer is the SiO2-rich surface layer.  0  50  100  150  200  250  300  350  400  SiO2 rich  ZrO2  t [hrs]  3  ZS15  ZS20    ZS30  ZS15  ZS20    ZS30  ZS15  ZS20  ZS30  4  8  d  [  µ  m  ]  Oxide thickness vs. time at 1550°C  Fig. 7.  The total oxide thickness of  the ZS15, ZS20, and ZS30 com posites tested at various dwelling times. Calculated by adding the mea sured thicknesses of  the SiO2—rich surface  layer and the underlying  ZrO2 layer.  484  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 92, No. 2    \\x0c', 'February 2009  Oxidation of ZrB2-SiC: Influence of SiC Content on Oxide Phase Formation  485  Fig. 8.  Calculated ternary phase diagram of an isothermal section of the ZrO2-SiO2-B2O3 system at 15001C.11 The annotations in the lower right shows the equilibrium compositions expected for the fresh products of oxidation of zirconium diboride with 15, 20, or 30 vol% SiC and in the upper left shows  the compositions anticipated after the B2O3 evaporates from the BSZ liquids.  composites with a higher SiC concentration, as on the ZS20 and  liquid and solid-phase assemblage where composites with more  ZS30 composites compared with the ZS15 composite.  SiC have BSZ liquid with a  larger  SiO2/B2O3 in a smaller cell density.  ratio, which  results  The cross-sectional analysis for all three composites showed  that a larger content of SiC yields a more oxidation-resistant  composite,  indicated by the decrease in total oxide scale thick dissolves  less ZrO2, and thus The influence of SiC content on the oxidation rate of ZrB2-SiC composites can be understood in terms of the composition and  ness with increasing SiC content. The increased oxidation resis amount of BSZ liquid.  tance of the higher SiC-containing materials could be explained  by the composition of the BSZ liquid phase. With more SiC in the  composite, there is less B2O3 content in the BSZ liquid. With less B2O3, the liquid is more viscous so that oxygen diffusion through the liquid is slower, which would result in less inward oxygen  diffusion through the middle of  the cell where the BSZ liquid  squeezes out (SiO2-rich BSZ liquid retards inward oxygen diffusion better than B2O3-rich liquid). Also, higher SiC-containing material will form a SiO2-rich surface layer quicker than a lower SiC-containing material, thus hindering oxygen diffusion inward  towards  the bulk material  sooner and thus more  effectively.  Also,  if a SiO2-rich surface layer suppress the evaporation of the B2O3 and a protective surface  is formed faster,  then it can  layer can be formed quicker.  V.  Conclusions  ZrB2-SiC composites with 20 and 30 vol% SiC oxidized at 15501C show convection cell features similar to convection cells  formed on an oxidized ZrB2-15 vol% SiC composite. Convection cell density decreases with increased oxidation time and SiC  concentration in ZrB2-SiC composites. The effect of SiC concentration on convection cell density can be correlated to the  Table I. The Amount of Precipitated Secondary ZrO2(s); BSZ Liquid Composition; and Estimated Viscosity,10,11 for the Three ZrB2-SiC Composites Estimated by Using the ZrO2- B2O3-SiO2 Ternary Phase Diagram Shown in Fig. 8  SiC concentration  (mol%)  BSZ composition  ZrO2, SiO2, B2O2  (mol%)  Precipitated  ZrO2(S) (mol% of  total composition)  15  20  30  11, 18, 71  8, 25, 67  6, 37, 57  34  23  12  Estimated  viscosity (Pa \\x01 s)  6.3 \\x02 102 5.0 \\x02 103 3.9 \\x02 104  BSZ, ZrO2-SiO2-B2O3.  Acknowledgments  We thank Dr. David Shifler of the Office of Naval Research for supporting the  research under  contract N00014-02-1-0034,  and Drs. Alida Bellosi, Frederick  Monteverde, Gregory Hilmas, and William Farhenholtz for providing the ZrB2- SiC composites used for this study and valuable discussions.  References  1M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurface: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 2M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey,  ‘‘Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and Zir conium Compounds,’’ J. Eur. Ceram. Soc., 19, 2405-14 (1999). 3R. Telle, L. S. Sigl, and K. Takagi,  ‘‘Transition Metal Boride Ceramics’’;  pp. 803-945 in Handbook of Ceramic Hard Materials, Vol. 2, Edited by R. Reidel.  Wiley-VCH, Weinheim, Germany, 2000. 4W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of SiC Addition on the  Oxidation of ZrB,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 5R. A. Cutler, ‘‘Engineering Properties of Borides’’; pp. 787-803 in Ceramics and  Glasses, Engineering Materials Handbook, Vol. 4, Edited by S. J. Schneider. ASM  International, Materials Park, OH, 1992. 6S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. 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Halloran,  ‘‘Formation of Oxide Films on ZrB2- SiC Composites During Oxidation: Relation of Subscale Recession to Liquid  Oxide Flow,’’ J. Am. Ceram. Soc., 91 [11] 3652-8 (2008). 13S. N. Karlsdottir, J. W. Halloran, F. Monteverde, and A. Bellosi,  ‘‘Oxidation  of ZrB2-SiC: Comparison of Furnace Heated Coupons and Self-Heated Ribbon Specimens’’; Proceedings of the International Conference on Advanced Ceramics and  Composites, Daytona Beach FL, January 21-26, (2007).  \\x0c', '486  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 92, No. 2  14Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby,  ‘‘Oxidation of  17A. Wackernagel, C. Massone, G. Hoefler, E. Steinbauer, H. Kerl, and P. Wolf,  ZrB2-SiC Ceramics under Atmospheric and Reentry Conditions,’’ Refractories Appl. Trans., 1 [2] 1-8 (2005). 15A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, ‘‘Oxidation of Zirconium Diboride-Silicon Carbide at 15001C at a Low Partial Pressure of Oxygen,’’ J. Am.  Ceram. Soc., 89 [10] 3240-5 (2006). 16S. N. Karlsdottir and J. W. Halloran,  ‘‘Formation of Oxide Films on ZrB2- SiC Composites During Oxidation: Evolution with Time and Temperature,’’  J. Am. Ceram. Soc., 2008, (accepted).  ‘‘Plasmacytoid Dendritic Cells are Absent in Skin Lesions of Polymorphic Light  Eruption,’’ Photodermatol., Photoimmunol. Photomed., 23, 24-8 (2007). 18H. Daims and M Wagner, ‘‘Quantification of Uncultured Microorganisms by  Fluorescence Microscopy and Digital Image Analysis,’’ Appl. Microbiol. Biotech nol., 75, 237-48 (2007). 19R. Jabra, J. Phalippau, and J. Zarzicki, Oxide Glasses by Hot-Pressing,’’ J. Non-crystalline Solids, 42, 489-98 (1980). &  ‘‘Synthesis of Binary Glass-Forming  \\x0c']"
},{
  "_id": 183,
  "PDF": "Oxidation of ZrB2-ZrO2 and ZrB2-ZrO2-SiC Ceramics.pdf",
  "Text": "['Advanced Materials Research Vol 66 (2009) pp 61-64 © (2009) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/AMR.66.61  Online: 2009-04-01  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, www.ttp.net. (ID: 128.6.218.72, Rutgers University Libraries, New Brunswick, USA-08/04/15,05:10:20)  Oxidation of ZrB2/ZrO2 and ZrB2/ZrO2/SiC Ceramics YUAN Huipinga, SONG Jianrong, LI Junguo, SHEN Qiang, ZHANG Lianmeng* State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, China *Corresponding Author, lmzhang@whut.edu.cn Keywords: ZrB2 composite, SPS, oxidation  Abstract. The purpose of this study was to investigate the oxidation of ZrB2/ZrO2 (ZZ) and ZrB2/ZrO2/SiC (ZZS) ceramics. The ceramics were fabricated by spark plasma sintering (SPS) at 1900°C and exposed to ten-minute oxidation cycles in stagnant air at 1200°C in a box furnace with molybdenum disilicide heating elements. Results of relative density, surface phase change and the rate of weight growth show that the addition of ZrO2 improved the sintering properties of ZrB2 ceramics. While the resistance to oxidation declined with the increase content of ZrO2. And the addition of SiC improved the resistance to oxidation of ZrB2/ZrO2 composites significantly. Introduction The borides, carbides, and nitrides of the early transition metals are considered Ultra-High Temperature Ceramics (UHTCs) because of melting temperatures above 3000 K, high hardness, and resistance to chemical attack [1-3]. Heating ZrB2 in air produces a scale composed of ZrO2 and B2O3 [4-8].William G. Fahrenholtz calculated a volatility diagram for temperatures of 1000, 1800, and 2500 K to understand the oxidation of ZrB2 [9]. E. Opila, S. Levine and J. Lorincz investigated the oxidation of several compositions of ZrB2and HfB2-based UHTCs in stagnant air at 1627°C in ten minute cycles for times up to 100 min [10]. The purpose of this work is to explore effects of ZrO2 and SiC addition to UHTCs in hopes of improving oxidation properties for this class of materials to enable their use in space transportation leading edge applications for short times at very high temperatures. This paper represents a preliminary investigation on the oxidation resistance of ZrB2/ZrO2 and ZrB2/ZrO2/SiC ceramics. Experimental Procedure Powder Processing. ZrB2-based UHTCs were prepared from ZrB2, ZrO2 and SiC powders. ZrB2 powder for research only was used in this study. ZrB2 powder had a purity of 99.5% (metal basis excluding Hf) and an average particle size of 44 µm. SiC powder had a purity of 99.99%  and an average particle size of 1.211.49 µm. Eight compositions, their designation, and processing history are summarized in Table 1. The powders were mixed and milled in an agate pestle for 4 h.  Table 1 The ingredient of ZZ and ZZS ceramics Samples Z95Z5 Z85Z15 Z75Z25 Z65Z35 Z80Z15S5 Z70Z15S15 Z60Z15S25 Z50Z15S35 ZrB2[vol%] 95 85 75 65 80 70 60 50 ZrO2[vol%] 5 15 25 35 15 15 15 15 SiC[vol%]    0 0 0 0 5 15 25 35  Sample Preparation and Sintering. Milled powders were sintered by Spark Plasma Sintering System using a graphite die. Powders were heated under vacuum (6 Pa) to 1900°C. A heating rate of 100°C/min was used from room temperature to 1900°C with hold times of 10 min. After the sintering, the furnace was cooled to room temperature naturally in 40 min.  \\x0c', '62  Advanced Synthesis and Processing Technology for Materials   Sample Characterization. Bulk density was determined using Archimedes’ technique with alcohol as the immersing medium. Relative density was calculated by driving the bulk density by the measured true density. X-ray diffraction analysis was performed to identify crystalline phases. Oxidation. Initial sample weights and surface crystalline phases were recorded. The ceramics were exposed to stagnant air at 1200°C in a box furnace with molybdenum disilicide heating elements for 10 min, and then took out and cooled in air to room temperature.  Sample weights were recorded again when cooled to room temperature. Repeat the 10-min oxidation cycles for 24 times.  Results and Discussion Sintering Behavior. The effect of ZrO2 and SiC as a second and third phase was evaluated with a series of densification experiments (Fig.1). The addition of 15, 25, or 35vol% ZrO2 increased the relative density of ZZ ceramics to 98.61%, 99.81% or 99.45%, respectively(solid square symbols in Fig. 1, (a)). The ZZS ceramics showed proximate relative density. When SiC was added to the ZxZ15 powder (open square symbols in Fig. 1, (b)), the pellets achieved a relative density between 98.86%~99.63% after sintering at 1900°C for 10 min. Based on these results, the addition of ZrO2 less than 25vol% led to more complete densification of ZrB2 at 1900°C. The content of SiC had no direct influence on the ZZS ceramics relative density.   510152025303593949596979899100  Relative Density/%Content of ZrO2/vol%ZxZy (a)  510152025303593949596979899100   Content of SiC/vol%Relative Density/% ZxZ15Sz(b) Fig. 1  Relative density as a function of the content of the (a) ZrO2 in ZxZy ceramics and (b) SiC in ZxZ15Sz ceramics  2030405060   ZrB2t-ZrO2m-ZrO2Z65Z35Z95Z52Theta/°Intensity Z85Z15  Z75Z25   20304050607080Z80Z15S5  ZrB2t-ZrO2SiCm-ZrO22Theta/°IntensityZ60Z15S25     Fig. 2. X-ray diffraction analysis of ZZ and ZZS ceramics (m)=monoclinic, (t)=tetragonal.  The crystalline phases of ZZ and ZZS ceramics after sintered at 1900°C by SPS were showed in Fig. 2. X-ray diffraction analysis show that m-ZrO2 can be found in ZZ and ZZS ceramics except ZrB2, and t-ZrO2. The m-ZrO2 existence has been attributed to phase change between t-ZrO2 and m-ZrO2. When cooled to 1117°C, t-ZrO2 will transfer to m-ZrO2.  Oxidation. Oxide formation is visible on the surface of ZZ ceramics tested at 1200°C. Loose gray formations was observed on samples exposed at 1200°C. With increasing exposure time, the oxide scale became white. The mass change reflects a combination of mass loss and mass gain. ZZ ceramics have almost same mass gain before 120min. Then Z75Z25 and Z65Z35 get a rapid, liner mass gain. XRD analyses show the surface oxidation product is largely monoclinic ZrO2 under this condition. \\x0c', 'Advanced Materials Research Vol. 66  63   0601201802400102030405060708090 Z65Z35 Z75Z25 Z85Z15 Z95Z5   Oxidation time(min)Mass change((((mg/cm2))))(a)  2025303540455055606560min    t-ZrO2  0min Intensity2Theta/°  (b) ZrB2 m-ZrO2 B2O310min   20304050607080 m-ZrO2    Z65Z35Z95Z5  Z85Z15  2Theta/° IntensityZ75Z25   (c) Fig. 3. The mass change of ZZ (a), X-ray diffraction pattern of Z95Z5 (b),  and the surface phase of ZZ ceramics exposed to air at 1200°C for 6 cycles (c).  ZrB2 + 5/2O2 → ZrO2 + B2O3(l)                                                                                                             (1) B2O3(l) → B2O3(g)                                                                                                                              (2) t-ZrO2→m-ZrO2                                                                                                                             (3)  The ZrB2 undergo stoichiometric oxidation when exposed to air at 1200°C at the surface. Reaction (1), (2)and (3) occurred at the surface. At the beginning the ZrO2 and B2O3 formed a continuous layer((b), Fig.3) and then only the ZrO2 remained due to B2O3 evaporation. The phase change between m-ZrO2 and t-ZrO2 (at about 1117°C) during the oxidation cycles result in the volume change and cracks which lead to oxidation. The B2O3 film inhibits further oxidation before 120min. With B2O3 evaporation after 120min, specimens began to gain mass greatly as the mass of oxide formed is greater than the mass of diboride reacted plus the mass of B2O3 lost.  06012018005101520 Mass change((((mg/cm2))))Oxidation time(min)Z50Z15S35Z60Z15S25Z70Z15S15Z80Z15S5(a) 20304050607080  0min2Theta/° Intensity  10min(b)  60minZrB2m-ZrO2t-ZrO2SiCZrSiO4   170min Fig. 4. The mass change of ZZS ceramics after 17 oxidation cycles in air at 1200°C  Oxide formation is visible on the surface of ZZS ceramics tested at 1200°C. Glass formation was observed on samples exposed at 1200°C. With increasing exposure time, the thickness of glass layer increased. Both ZrB2 and SiC oxidation occurred at the surface of ZZS ceramics. (Shown in Fig.4) Before 10min a similar response that the ZrO2 and B2O3 formation result in mass gain has been observed for ZZS ceramics. Z80Z15S5 ceramics get a linear mass gain after 20min. Other ZZS specimens show little mass gain (less than 8mg/cm2). The mass and surface phase of Z60Z15Z25 almost did not change after 10min. As shown in Fig. 4, ZrO2 and ZrSiO4 are mainly oxidation products.  Similar with ZZ ceramics, the B2O3 film inhibits further oxidation at the beginning. With oxidation time increasing, the ZrSiO4 (Reaction (4) and (5)) inhibit further oxidation in ZZS ceramics. Its rate limiting step for oxidation is the B2O3 evaporation from the matrix and the transport of oxygen through B2O3, which result in parabolic (diffusion-limited growth) kinetics for mass gain and the oxide layer thickness. SiC + 3/2O2 = SiO2 + CO(g)                                                                                                             (4) SiO2 + ZrO2 = ZrSiO4                                                                                                                        (5)  \\x0c', '64  Advanced Synthesis and Processing Technology for Materials       Fig. 5. Electron probe microanalysis showing the layered structure that develops on ZZS that was exposed to air at 1200°C for 17 cycles (170 min) A photomicrograph of the cross-section of ZZS after oxidation for seventeen 10-min cycle at 1200°C is shown in Fig. 5. The gray crystal is ZrB2 and lighter colored is ZrO2 phase. The dark phase is SiC, which is distributed in ZrB2/ZrO2 background. Oxidation of ZZS at 1200°C for 170 min in air produced a structure that consisted of three layers (shown in Fig. 5): (1) a continuous surface oxide layer consist of ZrO2 and ZrSiO4 (about 8µm thicknesses), (2) a SiC-depleted layer (note that SiC is only partially depleted and not fully removed) that may be ZrO2 and ZrB2 (about 20µm thicknesses), and (3) unaffected ZrB2/ZrO2/SiC. Fine ZrO2 and ZrSiO4 grains less than 3µm are the mainly oxidation products in the oxide layer. The formation of this layer structure is consistent with observations of other investigators who have studied ZrB2-SiC oxidation [11, 12]. Summary Dense ZrB2/ZrO2 and ZrB2/ZrO2/SiC ceramics can be prepared at temperatures as low as 1900°C by SPS. The addition of ZrO2 is not protective and SiC which alter to ZrSiO4 after oxidation is protective for ZZ and ZZS ceramics evolve 10min oxidation cycles at 1200°C as the formation of condensed ZrSiO4 film limit B2O3 evaporation from the matrix and the transport of oxygen through B2O3. Acknowledgement This work is supported by Program for Innovation Team in Hubei Province No. 2008CDA011.  References [1] R. Telle, L.W. Sigl and K. Takagi, Boride-Based Hard Mater, in: Handbook of Ceramic Hard Materials, edited by R. Riedel. Wiley-VHC, Weinheim, (2000). [2] T. A. Jackson, D. R. Eklund and A. J. Fink: J. Mater. Sci. Vol. 39 (2004), p. 5905 [3] D.M.V. Wie, D.G. Drewry Jr., D.E. King and C.M. Hudson: J. Mater. Sci. Vol. 39 (2004), p. 5915 [4] W.C. Tripp and H.C. Graham: J. Electrochem. Soc. Vol. 118 (1971), p. 1195 [5] J.B. Berkowitz-Mattuck: J. Electrochem. Soc. Vol. 113 (1966), p. 908 [6] R.J. Irving and I.G. Worsley: J. Less-Common Metals Vol. 16 (1968), p. 103 [7] I.G. Talmy, J.A. Zaykoski and M.M. Opeka: Ceram. Eng. Sci. Proc. Vol. 19 (1998), p. 105 [8] A.K. Kuriakose and J.L. Margrave: J. Electrochem. Soc. Vol. 111 (1964), p. 827 [9] W. G. Fahrenholtz: J. Am. Ceram. Soc. Vol.88 (2005), p. 3509 [10] E. Opila, S. Levine and J. Lorinc: J. Mater. Sci. Vol. 39 (2004), p. 5969 [11] W. G. Fahrenholtz: J. Am. Ceram. Soc., Vol. 90 (2007), p. 143 [12] S.N. Karlsdottir and J.W. Halloran: J. Am. Ceram. Soc. Vol. 90 (2008), p. 272 ZrSiO4+ZrO2 SiC-depleted layer ZrB2/ZrO2/SiC \\x0c', 'Advanced Synthesis and Processing Technology for Materials   10.4028/www.scientific.net/AMR.66   Oxidation of ZrB2/ZrO2 and ZrB2/ZrO2/SiC Ceramics   10.4028/www.scientific.net/AMR.66.61   DOI References  [3] D.M.V. Wie, D.G. Drewry Jr., D.E. King and C.M. Hudson: J. Mater. Sci. Vol. 39 (2004), p. 5915  doi:10.1023/B:JMSC.0000041688.68135.8b  [2] T. A. Jackson, D. R. Eklund and A. J. Fink: J. Mater. Sci. Vol. 39 (2004), p. 5905  doi:10.1023/B:JMSC.0000041687.37448.06  [4] W.C. Tripp and H.C. Graham: J. Electrochem. Soc. Vol. 118 (1971), p. 1195  doi:10.1149/1.2408279  [5] J.B. Berkowitz-Mattuck: J. Electrochem. Soc. Vol. 113 (1966), p. 908  doi:10.1149/1.2424154  [6] R.J. Irving and I.G. Worsley: J. Less-Common Metals Vol. 16 (1968), p. 103  doi:10.1016/0022-5088(68)90067-2  [8] A.K. Kuriakose and J.L. Margrave: J. Electrochem. Soc. Vol. 111 (1964), p. 827  doi:10.1149/1.2426263  [10] E. Opila, S. Levine and J. Lorinc: J. Mater. Sci. Vol. 39 (2004), p. 5969  doi:10.1023/B:JMSC.0000041693.32531.d1  [11] W. G. Fahrenholtz: J. Am. Ceram. Soc., Vol. 90 (2007), p. 143  doi:10.1111/j.1551-2916.2006.01329.x  [1] R. Telle, L.W. Sigl and K. Takagi, Boride-Based Hard Mater, in: Handbook of Ceramic Hard Materials,  edited by R. Riedel. Wiley-VHC, Weinheim, (2000).  doi:10.1002/9783527618217.ch22         \\x0c']"
},{
  "_id": 184,
  "PDF": "Oxidation of ZrB2–SiC Ultrahigh-Temperature Ceramic Composites in Dissociated Air.pdf",
  "Text": "['JOURNAL OF THERMOPHYSICS AND HEAT TRANSFER  Vol. 23, No. 2, April-June 2009  Oxidation of ZrB2-SiC Ultrahigh-Temperature Ceramic Composites in Dissociated Air  Jochen Marschall∗ and Dušan A. Pejaković†  SRI International, Menlo Park, California 94025  William G. Fahrenholtz,‡ Greg E. Hilmas,§ and Sumin Zhu¶  Missouri University of Science and Technology, Rolla, Missouri 65409  Jerry Ridge∗∗  NASA Ames Research Center, Moffett Field, California 94035  and Douglas G. Fletcher,†† Cem O. Asma,‡‡ and Jan Thömel§§  von Kármán Institute for Fluid Dynamics, 1640 Rhode-Saint-Genèse, Belgium  DOI: 10.2514/1.39970  The  oxidation  behavior  and surface  properties  of  hot-pressed ZrB2 -SiC ultrahigh-temperature  ceramic  composites are investigated under aerothermal heating conditions in the high-temperature, low-pressure partially  dissociated airstream of the 1.2 MW Plasmatron facility at the von Kármán Institute for Fluid Dynamics. Samples  are oxidized at different flow enthalpies for exposure times of up to 20 min at surface temperatures ranging from 1250 to 1575\\x0eC. The microstructure and composition of the resulting oxide layers are characterized using electron and  optical microscopies, x-ray diffraction, and energy-dispersive x-ray analysis. Comparisons are made with samples  oxidized under similar temperature and pressure conditions in a furnace test environment in which atomic oxygen  concentrations are negligible. Changes  in surface optical properties are documented using spectral reflectance  measurements, and effective catalytic efficiencies are estimated using computational fluid dynamics calculations  together with measured surface temperatures and heat fluxes.  H  _m  P  q  Rm T u1 , u2 u, v  = = = = = = = =  Nomenclature  enthalpy, J \\x01 kg\\x001 \\x01 K\\x001 mass flow rate, kg \\x01 s\\x001 pressure, Pa heat flux, W \\x01 m\\x002 model radius, m temperature, K dimensionless velocity gradients in the x and y directions velocity in the x and y directions, m \\x01 s\\x001  V x, y  \\r 0  \\x01  \\x0e  \"  \\x1b  = = = = = = = =  dimensionless velocity in the y direction surface parallel and normal coordinate axes atomic recombination coefficient catalytic efficiency nondimensional boundary-layer thickness boundary-layer thickness, m emittance Stefan-Boltzmann constant;  5:67 \\x02 10\\x008 W \\x01 m\\x002 \\x01 K\\x004  Received 23 July 2008; revision received 16 January 2009; accepted for publication 18 January 2009. Copyright © 2009 by the American Institute of Aeronautics and Astronautics, Inc. The U.S. Government has a royalty-free license to exercise all rights under the copyright claimed herein for Governmental purposes. All other rights are reserved by the copyright owner. Copies of this paper may be made for personal or internal use, on condition that the copier pay the $10.00 per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923; include the code 0887-8722/09 $10.00 in correspondence with the CCC. ∗Senior Research Scientist, Molecular Physics Laboratory, 333 Ravenswood Avenue; jochen.marschall@sri.com. Senior Member AIAA. †Research Physicist, Molecular Physics Laboratory, 333 Ravenswood Avenue. Member AIAA. ‡Professor, Department of Materials Science and Engineering, 223 McNutt Hall; billf@mst.edu. §Professor, Department of Materials Science and Engineering, 223 McNutt Hall; ghilmas@mst.edu. ¶Ph.D. Candidate, Department of Materials Science and Engineering, 223 McNutt Hall; Sumin.Zhu@us.vesuvius.com. ∗∗Research Engineer, Thermal Protection Materials and Systems Branch, Building 234; also, ELORET Corporation, Sunnyvale, California 94086; jerome.w.ridge@nasa.gov. ††Past Head, Aeronautics and Aerospace Department; currently, Professor, Mechanical Engineering, University of Vermont, 201 Votey Hall, 33 Colchester Avenue, Burlington, VT, 05405; douglas.fletcher@uvm.edu. Associate Fellow AIAA. ‡‡Research Engineer. Aeronautics and Aerospace Department, Ch. De Waterloo 72; asma@vki.ac.be. §§Ph.D. Candidate, Aeronautics and Aerospace Department. Ch. De Waterloo 72; jan.thoemel@vki.ac.be. Student Member AIAA.  Subscripts  cond cw dyn  e  stat  w  = = = = = =  conduction cold wall dynamic boundary-layer edge static wall  I.  Introduction  U LTRAHIGH-TEMPERATURE ceramic (UHTC) composites containing transition metal borides, carbides, and nitrides are under intensive investigation for applications in extreme environments [1,2]. UHTC composites based on ZrB2 and HfB2 , in combination with silica formers such as SiC and MoSi2 , are particularly attractive for high-temperature aerospace applications, such as leading-edge and control surface components on hypersonic vehicles. Both ZrB2 and HfB2 and their transition metal oxides, ZrO2 and HfO2 , have extremely high melting points (>2500\\x0eC). Hafnium and zirconium diborides have higher thermal conductivities compared with their carbide and nitride analogs, giving them a performance advantage for sharp leading-edge components for which drawing heat away from the stagnation point region is essential. Composites based on ZrB2 have the additional benefits of lower density and lower cost compared with HfB2 -based composites. The oxidation behavior of ZrB2 and its composites has been studied by many researchers using thermal gravimetric analysis  267  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\r \\x0c', '268  MARSCHALL ET AL.  of ZrB2  (TGA) or conventional furnace environments [3-14]. During temperatures between 1200 and 1600\\x0eC isothermal oxidation at and pressures near 1 atm, an oxide scale consisting of two distinct layers develops on the surface of ZrB2 -SiC composites [4,12]. The outer layer is a dense oxide scale consisting of a silica-rich glass with some embedded ZrO2 crystallites. This glassy scale is separated from the underlying virgin material by a partially oxidized and somewhat porous interlayer, containing ZrO2 and/or ZrB2 , from which the SiC has been partially or fully depleted by active oxidation. Virgin ZrB2 -SiC surfaces oxidize initially through the parallel net  reactions ZrB2 \\x87 5=2O2 ! ZrO2 \\x87 B2O3 \\x85l\\x86 and SiC \\x87 3=2O2 !  SiO2 \\x87 CO\\x85g\\x86. Liquid B2O3 mixes with amorphous SiO2 to form a borosilicate glass that seals the surface. However, B2O3 has a high vapor pressure that causes it to evaporate preferentially from the glassy phase and leave behind a silica-rich material. As the silicarich-layer thickens, it limits the inward diffusion of oxygen to the virgin material below, slowing oxidation and lowering the oxygen partial pressure in the reaction zone [12]. At sufficiently low oxygen pressures, the oxidation becomes negligible and the oxidation of silicon carbide becomes “active,” proceeding by the reaction SiC \\x87 O2 ! SiO\\x85g\\x86 and CO\\x85g\\x86 [12,15]. Oxidation rates are commonly quantified by measuring the thickness of the oxide layers or by measuring the mass change of the specimen. For oxidation of ZrB2 these rates are normally proportional to the square root of time (i.e., parabolic kinetics) consistent with diffusion-limited oxidation. Oxidation under lower total pressures and higher temperatures can result in evaporation of the SiO2 in addition to the B2O3 , leaving behind a single outer layer of porous ZrO2 and resulting in rapid, linear oxidation kinetics [13]. In both conventional TGA and furnace experiments, the gas environment is in thermal equilibrium with the test specimen. The thermochemical makeup of the test gas is determined by the feed gas composition and test temperature and pressure through equilibrium chemistry. In contrast, in hypersonic flight environments, large temperature gradients exist between the leading-edge surfaces and the boundary layer and shock layer edges. The gas temperature in a shock layer can easily exceed 10,000\\x0eC, resulting in an energetic gas mixture whose thermochemical makeup may contain ions, atoms, and molecules in highly energetic states. The shocked gas will undergo thermochemical relaxation as it nears the colder surface, but typically will not reach chemical equilibrium at the temperature of the surface. One ramification of this thermochemical nonequilibrium is that UHTC materials on leading-edge surfaces may interact with significant concentrations of highly reactive atomic oxygen. In this paper, we report the results of oxidation tests performed in high-temperature, low-pressure dissociated airflows on hot-pressed composites containing 30 vol% SiC (ZrB2 specimens were manufactured at the Missouri University of Science and Technology (Missouri S&T) in Rolla, Missouri, and tested in the 1.2 MW Plasmatron facility of the von Kármán Institute for Fluid Dynamics (VKI) in Rhode-Saint-Genèse, Belgium. The microstructure and composition of the resulting oxide layers are characterized and compared with oxidation studies performed at SRI International in a furnace under similar temperature and total pressure conditions. Results are discussed in terms of thermodynamic considerations, as well as previously published  -30SiC). The  ZrB2 ZrB2  -SiC,  -30SiC  oxidation studies in furnace and arcjet environments. Changes in optical properties resulting from surface oxidation are characterized by spectral reflectance measurements. The effective catalytic efficiencies for oxygenand nitrogen-atom recombination on the oxidizing surfaces are estimated using computational fluid dynamics (CFD) simulations in conjunction with measured surface temperatures and heat fluxes.  II.  Materials  \\x1820\\x0eC min\\x001  UHTC specimens were prepared from commercial powders supplied by H.C. Starck (grade B ZrB2 and UF-10 SiC). Powders were attrition milled using tungsten carbide media for 2 h in hexane and then dried by rotary evaporation. The dried powders were put into graphite dies that were lined with boron-nitride-coated graphite foil. Densification was accomplished by hot pressing (Thermal Technology FP-20-3560). The hot press was heated under vacuum (\\x1820 Pa) at \\x1820\\x0eC \\x01 min\\x001 rough to 1450\\x0eC, held for 1 h, then heated at the same rate to 1650\\x0eC and held for another hour. The hot press was then backfilled to 1 atm with argon and heated at temperature of 1900\\x0eC, to a at which a uniaxial pressure of 32 MPa was applied and the specimens were held for 45 min. At the end of this hold time, the hot press was cooled at \\x1820\\x0eC \\x01 min\\x001 to room temperature and, at \\x181750\\x0eC, the load was removed. After removal from the dies, the hot-pressed samples were approximately 3-4 mm thick with diameters of about 32 mm. The large faces of the disks were diamond ground flat and parallel at Missouri S&T and then shipped to Bomas Machine Specialties, Inc. (Sommerville, Massachusetts), where a 30 deg bevel was diamond ground into the disk edge, leaving a small face with a diameter of 26.6 mm, as shown in Fig. 1a. This bevel was used to hold the samples in the standard ESA 50 mm stagnation point test fixture used in VKI’s Plasmatron facility; see Fig. 1b. Machined UHTC specimens were cleaned with acetone in an rinsed with distilled water, and dried in a 100\\x0eC ultrasonic bath, oven. The mass and thickness of each specimen were measured before testing. Specimen densities calculated from these measurements and the known sample geometry ranged from 5.22 to theoretical density of ZrB2 -30SiC, assuming is 5:23 g \\x01 cm\\x003 . The slightly higher specimen densities were due densities of and for SiC, to the presence of a small volume fraction of tungsten carbide contamination from the milling media [16].  5:36 g \\x01 cm\\x003 . The  3:21 g \\x01 cm\\x003  6:09 g \\x01 cm\\x003  for ZrB2  III.  Experiment  A.  Plasmatron Facility  Plasma oxidation experiments were performed in the 1.2 MW Plasmatron Facility at VKI [17,18]. This facility uses a highfrequency, high-power, high-voltage (400 kHz, 1.2 MW, 2 kV) solid-state power supply to generate a high-enthalpy gas flow by inductive coupling in a 160-mm-diam plasma torch. The plasma flow was directed into a 2.5-m-long, 1.4-m-diam vacuum chamber equipped with multiple portholes and windows for optical diagnostics. Test models and probes for measuring dynamic pressure  Water-Cooled Sting Arm Water-Cooled Sting Arm Water-Cooled Sting Arm  Ceramic Ceramic Ceramic Insulation Insulation Insulation  UHTC UHTC UHTC Sample Sample Sample  SiC Cover SiC Cover SiC Cover  L = ~3 mm L = ~3 mm L = ~3 mm  Ds = 26.6 mm Ds = 26.6 mm Ds = 26.6 mm  φ = 60 deg φ = 60 deg φ = 60 deg  a)  b)  Fig. 1  Geometry: a) UHTC test specimens, and b) stagnation point specimen holder.  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\x0c', 'MARSCHALL ET AL.  269  and stagnation point heating rates were mounted on water-cooled arms that could be swung into and out of the flow on demand. The cold-wall stagnation point heat flux, qcw , was measured during each experimental run using a water-cooled copper calorimeter installed flush with the surface along the centerline of the watercooled copper probe. Heat flux values were determined from the temperature change and mass flow rate of the water used to cool the calorimeter. The size and external shape of the probe was identical to the specimen holder so that important flow parameters (dynamic pressure, velocity gradient) were reproduced. Both the probe and calorimeter had polished copper surfaces that were assumed to be fully catalytic to Nand O-atom surface recombination. Sample surface temperatures were measured by a two-color pyrometer (Marathon Series MR1SC, Raytek Corporation, Santa Cruz, California) at an acquisition rate of 1 Hz. The pyrometer uses two overlapping infrared wavelength bands at 0.8-1.1 and 1:0-1:1 \\x16m to measure temperatures from 1000 to 3000\\x0eC. Thermal emission from the sample was collected through a glass window at an incident angle of 29 deg off normal. The pyrometer was previously calibrated with the same window in place using a blackbody radiation source; temperature measurements are believed to be accurate to \\x0610\\x0eC. Unlike arcjet wind tunnels, which operate in the supersonic flow regime, inductively coupled plasma (ICP) facilities like the Plasmatron at VKI are typically operated in the subsonic flow regime. For materials-science-oriented studies, this operating regime has the advantage that the conditions in the plasma stream are much more likely to be at or near thermal and chemical equilibrium, making estimates of the gas chemical composition at the specimen surface more robust. In addition, metal contaminates that are often present in arcjet flows because of the electrode sputtering are absent in ICP facilities because the plasma flows see no metal surfaces inside the torch head.  B.  Plasmatron Test Conditions  The critical environmental parameters driving UHTC oxidation are the sample temperature, the gas pressure, and the chemical composition of the gas interacting with the sample surface. These parameters are essential for correlating the extent of UHTC oxidation with different test environments and understanding the kinetic aspects of oxide scale formation. Whereas surface temperature and gas pressure can be measured directly, the gas composition at the specimen surface cannot. Computational fluid dynamics simulations of the plasma freestream and the boundary layer around the model holder are used to compute the gas composition from known Plasmatron operating parameters and available surface temperature, pressure, and heat flux measurements made on the sample and probe during each test run. This is a two-step process. First, the freestream plasma flow conditions at the boundary-layer edge along the stagnation point centerline are rebuilt for each Plasmatron operating condition, using measured quantities and known facility and geometric parameters; then, these boundary-layer-edge conditions are used together with the measured sample surface temperature to compute the gas composition and temperature through the boundary layer to the sample surface. This procedure also couples the measured surface temperature to the boundary-layer-edge conditions through an energy balance at the sample surface, providing an estimation of the catalytic efficiency of the oxidizing surface for Oand N-atom recombination. The CFD codes used for this procedure were developed at VKI and include the VKI boundary-layer code [19,20] and the VKI ICP code [21,22], both of which use the PÉGASE library to perform thermodynamic and transport property calculations [23]. The performance and validation of these codes and examples of their use in simulating flows in the Plasmatron can be found in the referenced papers. The ICP code simulates the flow inside the plasma torch and around the test sample in the vacuum chamber by solving the timeaveraged magnetohydrodynamic equation at low Mach and low magnetic Reynolds numbers, assuming axisymmetric flow and local thermodynamic equilibrium. The boundary-layer code solves the boundary-layer equations for an axisymmetric or two-dimensional  steady laminar flow of chemically reacting gas over a catalytic surface, including thermal and chemical nonequilibrium. The code provides computations of the stagnation point heat flux with a functional dependence:  qw \\x88 qw \\x85\\r 0 ; Tw ; Pe ; Te ; \\x01; u1 ; u2 ; Ve \\x86  (1)  u1  where \\r 0 is the catalytic recombination efficiency, Tw is the wall temperature, Te and Pe are the gas temperature and pressure at the boundary-layer edge, \\x01 is the nondimensional boundary-layer thickness, Ve is the nondimensional axial velocity, is the nondimensional radial velocity gradient, and u2 is the nondimensional axial derivative of the radial velocity gradient. Figure 2 illustrates the relationship of the stagnation point boundary-layer region to the various parameters given earlier; u and v are the velocities along the x and y axes, and Rm is the radius of the axisymmetric test model. Together, these two codes were used to rebuild the flow conditions at the boundary-layer edge from the cold-wall heat flux and dynamic pressure measurements. The four nondimensional boundary-layer parameters were computed by the ICP code for each Plasmatron test condition, given the test and Plasmatron geometries, the power coupled into the gas, the gas mass flow, and the static pressure. With the assumption that the local thermodynamic equilibrium holds at the boundary-layer edge, these computed nondimensional parameters served as inputs to the boundary-layer code. By taking the measured cold-wall heat flux as fully catalytic (Tw \\x88 300 K, \\r 0 \\x88 1), values of Te and Ve could be iteratively adjusted until the computed and measured heat fluxes agreed. The final temperature, Te , determined the enthalpy of the gas at the boundary-layer edge, He . Once the boundary-layer edge conditions for a particular test condition were determined, the VKI boundary-layer code could qw \\x88 qw \\x85\\r 0 ; Tw \\x86, where compute a heat flux abacus, all other variables in Eq. (1) were fixed. The catalytic efficiency of a test specimen exposed to the same plasma flow was estimated by locating measured specimen surface temperature and hot-wall heat flux values on this abacus [24]. For each point on the abacus, the VKI boundary-layer code concurrently computed the concentrations of O, O2 , N, N2 , and NO species in the gas phase above the specimen surface. Thus, the location of each test on the heat flux abacus also determined the particular CFD solution used to estimate the gas composition at the sample surface. Note that the catalytic recombination efficiency \\r 0 used in the VKI boundary-layer code is the product of the species recombination coefficient, \\r (the fraction of species-wall collisions that results in recombination), and the energy accommodation coefficient, \\x0c (the fraction of exothermic recombination energy transferred to the wall). The distinction between these two recombination efficiencies has been discussed by Bedra and Balat-Pichelin [25]; the two efficiencies \\x0c \\x88 1. The VKI boundary-layer code also assumes that the catalytic are only equivalent in the case of ideal energy accommodation, efficiency is the same for O and N atoms (\\r 0 N ) and does not include NO as a heterogeneous reaction product.  O \\x88 \\r 0  Fig. 2  General schematic diagram of the stagnation point boundary layer region.  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\x0c', '270  MARSCHALL ET AL.  C.  Furnace Oxidations  Furnace oxidations were performed in a tube furnace with resistive molybdenum disilicide heating elements capable of operating at 1600\\x0eC in air (Model 1610-20, CM Furnaces, Bloomfield, New Jersey). Samples were place on top of a “D-shaped” alumina substrate and positioned at the furnace midpoint inside a 45-mmdiam high-purity alumina tube. The furnace tube was sealed with vacuum-tight, water-cooled compression fittings at both ends. Air (Matheson Ultra Zero Grade, 99.999%) and argon (Matheson Ultra High Purity Grade, 99.999%) were introduced at one end through electronic mass flow controllers at a rate of 1 l= min and evacuated at the opposite end by a high-capacity roots blower pumping system. Gas pressures were measured using a 100 torr capacitance manometer and maintained at the desired 104 Pa (75 torr) pressure level by adjusting a throttling valve inline with the vacuum system. 1575\\x0eC at Samples were heated to either 1250 or a rate of approximately 10\\x0eC \\x01 min\\x001 under a flow of pure argon. At the target temperature, the flow was switched from argon to air for the desired oxidation time. After the oxidation was completed, the gas flow was switched back to argon as the sample was cooled at a similar rate. The UHTC samples exposed to representative heating cycles of 1250 or 1575\\x0eC under pure argon flows showed no surface oxidation, confirming that oxygen impurities in argon and air leaks into the vacuum system were insignificant.  D.  Sample Characterization  (2:5-18:0 \\x16m) was  Sample masses and thicknesses were measured before and after testing using a Mettler Toledo XP105 analytical balance with 0.01 mg resolution and a Mitutoyo Model 543-272 Digimatic Indicator with 0.01 mm resolution. Sample thickness was measured at five locations, at the center and at four symmetric locations on a concentric 1.0-cm-diam circle, and averaged. The room-temperature, hemispherical-directional spectral reflectance of several preand posttest UHTC samples was measured using two spectrophotometers fit with integrating spheres. A PerkinElmer Lambda 9 spectrophotometer (0:25-2:5 \\x16m) was used to span the near ultraviolet to near infrared range and a BIORAD FTS-40 spectrophotometer used for longer infrared wavelengths. Posttest samples were prepared for microscopy by cutting specimens perpendicular to their oxidized faces and polishing the cross sections to a using diamond abrasives. Microstructure and composition analyses were performed by field emission scanning electron microscopy (SEM, Model S4700, Hitachi, Japan) along with energy-dispersive x-ray spectroscopy (EDS, Model Phoenix, EDAX, Inc., Mahwah, New Jersey). Grazing incidence x-ray diffraction (GXRD; X’Pert MRD, Panalytical, Almelo, The Netherlands) with Ni-filtered Cu K\\x0b radiation was used to determine the crystalline phases present in the oxidized surface in a penetration depth of less than \\x18200 nm into the specimen. To layers. The grazing angle for GXRD was set to 1 deg, which resulted examine the partially SiC-depleted interlayer with this technique, the silica-rich outer layer was removed by polishing parallel to the original surface. An optical microscope was used to verify material  0:25 \\x16m finish  removal so that the partially SiC-depleted interlayer was fully exposed. X-ray photoelectron spectroscopy (XPS; AXIS 165, Cratos Analytical, Kyoto, Japan) was used to acquire the nitrogen bonding information on the posttest surface layer of one sample (sample 9).  IV.  Experimental Results  A.  Measured and Derived Test Conditions  All UHTC tests reported here were performed with an air mass _m, of 16 g \\x01 s\\x001 and a static chamber pressure, Pstat , of flow rate, 104 Pa. Freestream enthalpy was varied by adjusting the Plasmatron power between 150 and 210 kW, resulting in measured dynamic pressures ranging from 24 to 39 Pa and cold-wall heat fluxes ranging from 51 to 123 W \\x01 cm\\x002 . The Plasmatron operating conditions and derived boundary-layer-edge conditions for each test are listed in Table 1. Preliminary tests showed that inserting and removing the UHTC specimens rapidly from the plasma stream under these conditions caused the SiC covers on the model holders to fracture due to thermal shock. Therefore, specimens were moved into the plasma flows at (\\x18100 kW) and the power was increased slowly to its low power target value. Similarly, after a desired exposure time, the power was decreased slowly to \\x18100 kW until the surface temperature dropped below 1100\\x0eC, before shutting the power off. Both the power rampup and ramp-down times were on the order of 1 min. Target test times were measured from the moment the sample was inserted into the flow until to the moment that the power was ramped down. The test times at quasi-steady surface temperatures were shorter and were estimated from the transient temperature profiles. A typical transient temperature profile is given in Fig. 3 for sample 9, tested at 210 kW for a target time of 10 min. The heating conditions for the entire UHTC test series are listed in Table 2. The second column of Table 2 gives both the target test time and (in brackets) the estimated time at a quasi-steady surface temperature. Quasi-steady surface temperatures typically exhibited fluctuations of tens of degrees Celsius (as seen in Fig. 3) and are reported to the nearest 10 deg in column 3 of Table 2. The hot-wall heat fluxes at the UHTC surfaces listed in Table 2 are not measured directly, but estimated from a radiative equilibrium calculation w \\x87 qcond , with \" \\x88 0:75 or 0.90, and qcond \\x88 0. The simplification qcond \\x88 0 causes a slight underestimation of the true hot-wall heat flux; however, this underestimation is small compared with errors introduced by uncertainty in the sample emittance. The emittance values \" \\x88 0:75 and 0.90 are reasonable for oxidizing ZrB2 -SiC surfaces, and each value is supported by some hot-wall heat flux estimated for \" \\x88 0:75 is 17% smaller indirect experimental evidence as will be discussed in Sec. IV.C. The than for \" \\x88 0:90, locating the tests at slightly different points \\x85Tw ; qw \\x86 on the heat flux abacus computed using the VKI boundary-layer code. This, in turn, leads to slightly different values of derived recombination coefficients and species concentrations at the specimen surface. Tables 3 and 4 list the values of \\r 0 ; the number densities of N2 , N, and O; and the number density ratios of O=O2 and O/NO for \" \\x88 0:75 and 0.90, respectively. The derived catalytic efficiencies are about 3- 7 times larger for \" \\x88 0:90 than for 0.75, because a greater chemical  as qw \\x88 \"\\x1bT 4  Table 1  Plasmatron operating conditions, where Pstat \\x88 104 Pa and _m \\x88 16 g \\x01 s\\x001 for all tests  Test-Sample  Power, kW  qcw , W \\x01 cm\\x002  Pdyn , Pa  He , kJ \\x01 g\\x001  Te , \\x0eC  ve , ms\\x001  5-20 6-16 7-15 8-14 9-13 10-12 11-11 12-5 13-6 14-10 15-9 16-4  150 160 160 160 170 170 170 160 190 190 210 210  51.0 71.5 73.0 75.5 90.5 87.5 90.0 71.0 109.5 105.0 122.5 116.5  24 25 26 26 28 27 27 24 31 31 38 39  11.2 15.2 15.4 15.9 18.6 17.8 18.7 15.3 21.9 21.0 24.5 21.9  4567 5164 5180 5229 5456 5394 5461 5169 5681 5623 5844 5684  87.5 98.2 100.4 101.4 110.0 106.7 108.2 96.3 121.6 120.0 126.0 136.4  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\x0c', 'MARSCHALL ET AL.  271  C  °  ,  E  R U  T  A  R  E P  M  E  T  E  C  A  F  R U  S  1700  1600  1500  1400  1300  1200  1100  1000  ~7.9 min at ~1570°C  Test Sample:15-9  0  2  4  6  8  10  12  14  TIME, min  Fig. 3  Transient surface temperature profile for test 15 sample 9 run at  210 kW for a target test time of 10 min.  contribution to heating is required to account for the higher estimated hot-wall heat flux. Figure 4 illustrates the uncertainty associated with the determination of catalytic efficiency, using case 10-12 as an example. The 1436\\x0eC, measured wall estimated hot-wall heat flux for \" \\x88 0:90 is 43:5 W \\x01 cm\\x002 ; temperature (to four digits) is and the this location is shown by the solid symbol in the heat flux abacus of Fig. 4, giving a derived catalytic efficiency of 2:65 \\x02 10\\x003 . Three sources contribute to the uncertainty of this value. First, the \\x0610C\\x0e uncertainty in the wall temperature measurement is shown by the horizontal error bars. Second, the freestream enthalpy can be determined reasonably within a 10% accuracy [26], limited mainly by the uncertainty in the cold-wall heat flux measurement. The contribution of this uncertainty is indicated by the pair of symmetric vertical error bars closest to the solid symbol. Finally, the more widely spread asymmetric vertical error bars indicate the uncertainty  Table 2  UHTC heating conditions  Test-Sample  Timea, min  Tw , \\x0eC  b, W \\x01 cm\\x002  \" \\x88 0:90  \" \\x88 0:75 qw  5-20 6-16 7-15 8-14 9-13 10-12 11-11 12-5 13-6 14-10 15-9 16-4  10 (8.2) 10 (7.7) 5 (3.2) 20 (18.0) 10 (8.3) 20 (18.3) 5(3.9) 16 (14.0) 20 (17.3) 10 (8.3) 10 (7.9)  15 \\x87 5  1240 1390 1360 1390 1460 1440 1460 1350 1510 1500 1570 1550  22.4 32.3 30.0 32.3 38.2 36.3 38.0 29.7 43.2 42.3 49.5 46.7  26.8 38.8 35.9 38.8 45.8 43.5 45.6 35.6 51.8 50.8 59.4 56.0  aTarget test time (approximate time at Tw ).  bFrom qw \\x88 \"\\x1bT 4 w \\x87 qcond , with qcond \\x88 0.  \\r 0  introduced by the surface temperature and emittance together into the radiative heat flux computation. The emittance uncertainty is as \\x870:1= \\x00 0:2, because by definition. The catalytic efficiency of 0:842 \\x02 10\\x003 derived by taken here emittance cannot exceed 1 assuming an emittance of 0.75 is also plotted in Fig. 4 as an open symbol and lies within this error bound. The emittance makes the dominant contribution to the radiative heat flux uncertainty and, thus, to the catalytic efficiency uncertainty. For this case, the uncertainty in is seen to approach an order of magnitude. Tables 3 and 4 show that atomic oxygen is the dominant oxygen species at the sample surface under all test conditions, with O2 and NO number densities typically 1 or more orders of magnitude lower than O-atom number densities. The O-atom number densities computed for \" \\x88 0:90 and 0.75 typically differ by less than 2%. Oxygen is essentially fully dissociated at all Plasmatron powers used in this study, but nitrogen dissociation is predicted to rise with increasing power. The Nand O-atom number densities at the sample surface become similar at the highest Plasmatron powers. Slightly larger N-atom concentrations near the surface are computed for \" \\x88 0:75 than for 0.90.  B.  Oxide Formation  All UHTC test specimens survived the Plasmatron exposures with minimal dimensional and mass changes and without any visual evidence of mechanical failure, such as chipping associated with spalling or specimen fractures induced by the thermal cycle between room temperature and the test temperature. This was true even during the initial exploratory tests that resulted in the failure of the SiC model cover. Figure 5 shows measured changes in the total sample mass and thickness as a function of target test time and surface temperature. For fixed test times, the samples generally gained the most mass and thicknesses at the lower temperatures, with these gains decreasing or turning into net losses at the higher temperatures. All posttest UHTC surfaces were black, shiny, and smooth on a macroscopic scale, indicative of the formation of a glassy oxide layer. Figure 6 shows SEM images of the oxidized surfaces (top row) and cross sections (bottom row) of samples 20, 16, and 9. These samples were exposed to the Plasmatron flow for target test times of 10 min at increasing power levels, attaining correspondingly higher surface temperatures. At this magnification, the surfaces are inhomogeneous, mostly covered by a glassy coating, but with regions of exposed crystalline material. Chemical analysis by EDS confirms the presence of Zr and O in the exposed crystallites and Si and O in the glassy phase, consistent with the anticipated oxidation products ZrO2 and SiO2 . Boron was not detected by EDS, but may be present in small quantities because light elements are difficult to quantify by EDS. The surface morphologies differ on the different samples. On sample 20, the glassy layer appears as a thin patchy film with many small, well-dispersed regions of exposed ZrO2 grains. On sample 16, most of the surface is covered with a uniform glassy film, with only a few exposed regions. On sample 9, the glass film appears thicker and  Table 3  Catalytic efficiencies and species number densities at the surface for \" \\x88 0:75  Test-Sample  \\r 0 , \\x0210\\x003  \\x89N2 \\x8a  5-20 6-16 7-15 8-14 9-13 10-12 11-11 12-5 13-6 14-10 15-9 16-4  0.333 1.04 0.285 0.681 0.936 0.842 0.911 0.349 0.680 0.828 1.13 1.15  2.97 2.44 2.37 2.33 2.04 2.13 2.04 2.40 1.70 1.79 1.52 1.69  Number Density, \\x021023 m\\x003  [N]  0.0444 0.379 0.516 0.504 0.728 0.666 0.734 0.488 1.05 0.955 1.16 1.01  [O]  1.75 1.53 1.56 1.52 1.41 1.44 1.41 1.56 1.31 1.33 1.23 1.28  \\x89O\\x8a=\\x89O2 \\x8a  [O]/[NO]  122 178 653 309 291 304 300 533 498 389 334 298  460 675 1420 973 933 959 949 1300 1220 1090 988 935  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970     \\x0c', 'more continuous than on sample 20, but with large partially exposed ZrO2 grains. The cross-sectional SEM images of Fig. 6 reveal a partially oxidized sublayer extending below the outer glassy layer in all three specimens. Figure 7 shows a higher-magnification SEM image of the two-layer oxide structure formed on sample 9, along with XRD spectra (plotted as diffracted x-ray intensity versus twice the diffraction angle) collected in each layer. The diffraction pattern suggests that monoclinic zirconia (m-ZrO2 ) is the major phase in the outer glassy layer (layer 1), but the amorphous glassy phases produce no diffraction peaks. The m-ZrO2 peaks must originate from exposed crystalline regions in layer 1. The SEM image reveals some porosity  regions of  in the sublayer (layer 2), and the XRD spectrum indicates m-ZrO2 , hexagonal \\x0b-SiC, and ZrB2 are present. The XRD spectrum of layer 1 shows an additional peak near 30 deg that was not detected in layer 2. This peak may indicate the presence of the high-temperature tetragonal phase of zirconia quenched by rapid cooling, or a partially nitrided zirconia phase, such as Zr7O11N2 [27]. Monteverde and Bellosi [6] cycled composites in atmospheric-pressure air up to 1350\\x0eC and assigned a similar peak in their surface XRD spectra to t-ZrO2 , hypothesizing that this phase was stabilized by impurities. The presence of a nitrided zirconia phase on the surface of sample 9 is plausible in view of the high N-atom concentration in the boundary layer during run 15; see Tables 3 and 4. The characteristic XRD peaks of Zr7O11N2 fit very well with the unidentified peaks shown in the layer 1 spectrum. Quantitative calculations using the layer 1 XRD spectrum and the structural parameters of m-ZrO2 and Zr7O11N2 single crystals indicate that the mass fraction of Zr7O11N2 is \\x1810%. The nitridation of ZrO2 on the surface on the sample 9 surface was additionally investigated by XPS. Figure 8 shows the N 1s core level feature in the measured XPS spectrum near 398 eV, confirming the presence of nitrogen atoms in the oxidized surface. Figures 9a and 9b show measurements of oxide layer thickness as \\x1810 measurements made a function of test temperature and time. Each point is the average of on high-magnification cross-sectional SEM images of regions with relatively uniform glassy surface oxide layers. The SEM images chosen for analysis, though necessarily selective, are representative of the oxide formed on each specimen. The error bars indicate scatter in the measurements for each analyzed image; larger variations in layer thicknesses are expected over the samples as a whole. The figures demonstrate that the glassy outer layer thickens with increasing temperature (1250, 1390, 1570\\x0eC) for a fixed test time of 10 min and thickens with increasing test times (5, 10, 20 min) in the 1360-1390\\x0eC range. Both at a fixed surface temperature the surface and sublayers in Fig. 9b exhibit parabolic growth in time, consistent with diffusion-limited oxidation. Figure 9a shows that growth of the glassy outer layer accelerates with increasing temperature, but that the sublayer is thickest at low temperatures, has a minimum at intermediate temperatures, and increases again at high temperatures. This behavior must reflect a balance between temperature-dependent transport and reaction kinetics, which increase with temperature and favor oxidation, and the thickening outer glassy layer, which hinders oxygen diffusion and suppresses oxidation. Figures 10a and 10b show surface SEM images of samples oxidized in a furnace for 20 min at 1250 and 1575\\x0eC, respectively, at a total air pressure of 104 Pa. The surface oxides formed on -30SiC specimens in the furnace environment and the Plasmatron environment under similar temperature and total are very different. At 1250\\x0eC, glassy oxide pressure conditions and exposed zirconia grains on the furnace-oxidized sample are segregated into much larger regions than on the Plasmatron-oxidized sample. Note the factor of 5 difference in magnification between the  that  t-ZrO2 ,  ZrB2  -SiC-Si3N4  the ZrB2  Table 4  Catalytic efficiencies and species number densities at the surface for \" \\x88 0:90  Number Density, \\x021023 m\\x003  Test-Sample  \\r 0 , \\x0210\\x003  \\x89N2 \\x8a  [N]  [O]  \\x89O\\x8a=\\x89O2 \\x8a  [O]/[NO]  5-20 6-16 7-15 8-14 9-13 10-12 11-11 12-5 13-6 14-10 15-9 16-4  2.33 2.91 1.78 2.36 2.81 2.65 2.78 1.87 2.56 2.73 3.27 3.19  3.05 2.66 2.57 2.55 2.30 2.35 2.27 2.61 1.93 2.02 1.77 1.92  0.000203 0.157 0.301 0.279 0.483 0.426 0.486 0.270 0.793 0.702 0.888 0.753  1.49 1.50 1.55 1.51 1.43 1.44 1.41 1.56 1.32 1.34 1.25 1.29  7.47 42.9 94.1 69.7 80.1 79.2 81.8 84.7 118 103 101 93.3  34.4 202 405 314 354 352 361 371 485 436 428 402  1420  1430  1440  1450  1460  30  35  40  45  50  55  ε = 0.90  ε = 0.75  γ\\', x10-3  4.6  3.6  2.65  1.6  0.6  0.265  q  w  ,  W  ⋅  c  m   2  Tw, °C  Fig. 4  Heat flux abacus for the determination of the catalytic efficiency  for  case 10-12 with uncertainty bars:  temperature  (horizontal),  free  stream enthalpy  (symmetric,  narrow vertical),  and  radiative  flux  (asymmetric, wide vertical).  0  5  10  15  20  25  1200  1250  1300  1350  1400  1450  1500  1550  1600  S  R U  F  A  C  E  T  E  M  E P  R  A  T  R U  E  ,  °  C  -19.7, -4  22.1, 2  35.5, 24  22.5, 2  14.7, -4  18.0, 2  -0.6, 2  4.1, -10  14.1, -6  21.4, 10  25.7, 8  TARGET TEST TIME, min  Fig. 5  Changes  in total  sample mass  (in milligrams)  and average  sample thickness (in microns) as a function of surface temperature and  target test time in the Plasmatron stream.  272  MARSCHALL ET AL.  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970         \\x0c', 'surface SEM images of Figs. 6 and 10. At the higher furnace temperature, the surface is made up of zirconia grains in a glassy matrix with no indication of a distinct overlaying glassy layer as observed for sample 9 oxidized in the Plasmatron. The furnaceoxidized sample at 1575\\x0eC appears white/gray in contrast to the black color of sample 9.  C.  Optical Properties  Changes in the hemispherical-directional spectral reflectance measured before and after exposure to the plasma flow are shown in  Fig. 11a for samples 5 and 6. Reflectance measurements performed on other preand posttest samples were very similar, with little discernible dependence on Plasmatron power or sample exposure time. In all cases, exposure to the plasma stream leads to a significant drop in reflectance, with posttest values around 0.1 for most of the spectral range. ZrB2 has metallic electrical properties, with a room-temperature electrical resistivity of about 8 \\x16  cm [28]. In the infrared region, a large drop in reflectance is consistent with the  Fig. 6  Posttest SEM images of samples 20, 16, and 9, which reached quasi-steady surface temperatures of 1240, 1390, and 1570\\x0eC during exposure to the  Plasmatron flow for 10 min at different power levels. The top row shows the oxidized surfaces, and the bottom row shows the cross sections showing  oxidized layers.  Fig. 7  XRD spectra of the oxidized layers on sample 9.  390             395             400             405             410  I  N  T  E  N  S  I  T  Y,  a  .  u  .  BINDING ENERGY, eV  N 1s  Fig. 8  The N 1s feature in the XPS spectrum of  the surface layer of  sample 9.  1250  1350  1450  1550  0  3  6  9  21  24  27  30  9  16  20  Layer 2  Layer 1  L  E Y A  R  T  H  I  C  K  N  S S E  ,  µ  m  TEMPERATURE, °C  0  5  10  15  20  0  3  10  15  20  25  30  14  16  15  Layer 2  Layer 1  L  E Y A  R  T  H  I  C  K  N  S S E  ,  µ  m  TARGET TEST TIME, min  a)  b)  Fig. 9  Oxide layer thicknesses estimated from SEM images of samples  similar times (\\x1810 min) at increasing temperatures, and b) samples 15, exposed to the Plasmatron stream: a) samples 20, 16, and 9 oxidized for (1360-1390\\x0eC)  16,  and  14  oxidized  at  similar  temperatures  for  increasing test times.  MARSCHALL ET AL.  273  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970           \\x0c', 'conversion of the UHTC surface layer from metallic to dielectric by oxidation of ZrB2 and SiC. In the visible region, this drop in reflectance explains the color change in specimen surfaces from a metallic gray to a deep black after exposure to the plasma flow. Some structure is evident in the measured spectra: a broad peak in the 10-13 \\x16m range for the pretest samples and a sharper feature centered near 9 \\x16m for the posttest samples. Figure 11b shows that similar features are found in theoretical reflectance spectra computed for a semi-infinite dielectric at normal incidence [29] using published optical constant data for SiC and SiO2 [30]. These reflectance features derive from the stretching mode vibrations of Si-C and Si-O bonds in the virgin and oxidized UHTC composite, respectively. Figure 11b includes a similar calculation for ZrO2 using the optical constants of Synowicki and Tiwald [31]. The predicted ZrO2 reflectance feature above 14 \\x16m is not observed in the experimental spectra. The presence of feature and the absence of the feature correlate reasonably well with the SEM microscopy of the posttest specimens that show a glassy outer layer with only minor amounts of exposed ZrO2 on the surface. Figures 12a and 12b show close-ups of the Si-O stretching mode reflectance feature for the same specimens in Figs. 9a and 9b, respectively. Growth of the reflectance feature in Fig. 12a correlates well with the significantly increasing glassy layer thickness with temperature in Fig. 9a (3:1-8:7 \\x16m), whereas Fig. 12b shows the relative insensitivity of the feature to more moderate thickness changes with time in Fig. 9b (2:5-4:3 \\x16m). Temperature-dependent total hemispherical emittance values can be estimated from room-temperature spectral reflectance measurements by using optical relations for an opaque solid and Kirchhoff’s law to convert spectral reflectance into spectral emittance and then averaging the spectral emittance weighted by the Planck function for different temperatures. This procedure presumes that weighting by the Planck function, rather than temperature-dependant optical constants, dominates the temperature-dependant total emittance. Figure 13 shows the results of such an estimate calculated from the spectral reflectance data of Fig. 11a. The predicted emittance of the oxidized specimens is quite high, near 0.9 over the entire temperature  the SiO2  ZrO2  range, whereas that of the virgin UHTC rises from a low value of \\x180:25 at room temperature to \\x180:6 at 2000\\x0eC. Because differences in the Plasmatron operating conditions had very little influence on reflectance spectra as a whole, the estimated emittances of all posttest  Fig. 10  Surface SEM images of samples oxidized in a furnace in 104 Pa air for 20 min at two different temperatures: a) 1250\\x0eC, and b) 1575\\x0eC.  0  3  6  9  12  15  18  0.0  0.2  0.4  0.6  0.8  1.0  Instrument Transition  Posttest  Pre-Test   Sample 5  Sample 6  R  E  F  L  E  C  T  A  C N  E  WAVELENGTH, µm  3  6  9  12  15  18  0.0  0.2  0.4  0.6  0.8  1.0  ZrO2  SiO2  SiC  R  E  F  L  E  C  T  A  C N  E  WAVELENGTH, µm  a)  b)  Fig. 11 Comparison of measured and calculated reflectance spectra: a) experimental hemispherical-directional spectral reflectance of samples 5 and 6 before and after exposure to the plasma flow, and b) normal spectral reflectance computed for SiO2 , SiC, and ZrO2 from published optical constants.  0  500  1000  1500  2000  0.0  0.2  0.4  0.6  0.8  1.0  Posttest  Pretest   Sample 5  Sample 6  E  M  I  T T  A  C N  E  TEMPERATURE, °C  Fig. 13  Temperature-dependent  total hemispherical  emittance  esti mated from room-temperature spectral reflectance measurements.  7  8  9  10  11  0.0  0.1  0.2  0.3   9  16  20  R  E  F  L  E  C  T  A  C N  E  WAVELENGTH, µm  7  8  9  10  11  0.0  0.1  0.2  0.3   15  16  14  R  E  F  L  E  C  T  A  C N  E  WAVELENGTH, µm  a)  b)  Fig. 12 Si-O stretching mode reflectance feature: a) samples 20, 16, and 9 tested for 10 min at 1250, 1390, and 1570\\x0eC; and b) samples 15, 16, and 14 tested at 1360-1390\\x0eC for 5, 10, and 20 min.  274  MARSCHALL ET AL.  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\x0c', 'MARSCHALL ET AL.  275  samples are essentially identical to those shown for samples 5 and 6 in Fig. 13. The emittance of a UHTC specimen will change as its surface progressively oxidizes during a Plasmatron test. The true instantaneous hemispherical emittance during a test run may lie between these preand posttest estimates, particularly during the initial heat-up of the specimen upon insertion. The lower value of \" \\x88 0:75 used in our hot-wall heat flux estimates lies within these bounds and is in good agreement with the measurements of total hemispherical emittance performed at temperatures between 1000 and 1700\\x0eC on ZrB2 -15SiC-2MoSi2 composites by Scatteia et al. [32] The higher value of \" \\x88 0:90 is consistent with our posttest reflectance measurements and also with the results of Monteverde and Savino [33], who recently tested hot-pressed ZrB2 80 kW plasma torch and derived an in situ emittance of \\x180:9 from simultaneous twoand one-color radiometry measurements.  -15SiC in an  D.  Catalytic Properties  In Fig. 14, we compare catalytic efficiencies derived in the present work with laboratory measurements reported by Scatteia et al. [32] and Marschall et al. [34] for similar hot-pressed ZrB2 -based composites. Measurements by Scatteia et al. were performed in the MESOX facility, which uses a solar furnace to heat specimens to high temperatures, a microwave-discharge flow tube to generate partially dissociated air, and an optical emission spectroscopy combined with actinometry to measure O-atom density gradients near the sample surface. Measurements by Marschall et al. were performed in a diffusion-tube side-arm reactor using a microwave discharge to generate binary O=O2 or N=N2 mixtures and twophoton laser-induced fluorescence detection of O-atom or N-atom concentration gradients. Experimental facilities, measurement techniques, and data analysis methods have been described in detail in the literature, for both the MESOX [35-39] and the side-arm reactor [40-44] setups. The MESOX and the side-arm reactor laboratory measurements derive values of the atom catalytic efficiency \\r , whereas the \\r 0 \\x88 \\r\\x0c. Marschall et al. performed measurements at relatively low experiments in the Plasmatron give values of the catalytic efficiency temperatures (up to 650\\x0eC). Silicon carbide oxidation is negligible in this temperature range. The primary oxidation products are B2O3 and ZrO2 , and the surface catalytic efficiency decreased once the melting point of B2O3 (\\x18450\\x0eC) was exceeded. All of the MESOX and Plasmatron data were obtained at temperatures well above 450\\x0eC. that \\r 0 Figure 14 shows values from Plasmatron experiments 1200-1700\\x0eC surface performed over a temperature range are factors of 5-20 times smaller (depending on the value of emittance used in the hot-wall heat flux estimate) than \\r values from the MESOX experiments. In this temperature range, SiC oxidation becomes increasingly important and, in both experiments, the oxidized UHTC surfaces are covered with a glassy silica-based oxide layer. If the silica-based surfaces are assumed to be similar in the two experiments, Fig. 14 implies that the exothermic recombination energy is not fully accommodated to the surface, with \\x0c roughly in  10-1  10-2  10-3  Y  C N  E  I  C  I  F F  E  C  I  T  Y  L  A  T  A  C  10-4 250  Present work,        ZrB2-30SiC,   γ\\'O = γ\\'N  Scatteia et al., [32]  ZrB 2-15SiC-2MoSi2               γ  O  Marschall et al., [34]    ZrB2-20SiC O, γ N    Before melting B2O3 O, γ N    After melting B2O3  γ  γ  ε = 0.90  ε = 0.75  500  750  1000 1250 1500 1750  TEMPERATURE, °  C  Fig. 14  Comparison of catalytic efficiencies determined for ZrB2 -SiC UHTC materials; measurements of Scatteia et al. [32] were performed in  the MESOX facility and those of Marschall et al.  [34]  in a side-arm  diffusion-tube reactor.  the 0.2-0.05 range. Rutigliano et al. [45] and Bedra et al. [46] have recently performed molecular dynamics simulations for N-atom recombination on \\x0c-crisobalite and O-atom recombination on quartz surfaces, respectively. In both studies, \\x0c was found to be about 0.4. Values of \\r around 0.01 are generally consistent with measurements on silica surfaces at these temperatures [25].  V.  Discussion  The oxide morphologies of the UHTC samples exposed to the Plasmatron stream suggest the following explanations. At temperatures below \\x181000\\x0eC, SiC oxidation is negligible and only the ZrB2 component of the UHTC composite oxidizes significantly [12]. The lowest-power Plasmatron test produces a slightly higher surface temperature of 1240\\x0eC (sample 20), at which SiC oxidation becomes more important, but is still slow compared with ZrB2 oxidation. This condition limits the formation of SiO2 relative to ZrO2 and B2O3 . However, the vapor pressure of liquid B2O3 at this temperature is already high, and the evaporation of boron oxide is significant. As a result, the borosilicate glass formed is only semiprotective as an oxygen barrier, and in-depth oxidation of ZrB2 and SiC proceeds readily, resulting in a relatively thick sublayer of partially oxidized UHTC material. At moderately higher surface temperatures, like the 1390\\x0eC experienced by sample 16, the oxidation rate of SiC becomes sufficiently high to form a continuous glassy layer over the oxidized UHTC surface, slowing in-depth oxidation by limiting inward oxygen diffusion. At the highest Plasmatron test temperature of 1570\\x0eC (sample 9), SiC oxidation is even more rapid and an even thicker silica-rich glassy layer develops. However, at this temperature the added protection provided by the thicker glassy layer is somewhat mitigated by faster oxygen diffusion rates in the glass, resulting in a partially oxidized sublayer thicker than sample 16, but still thinner than sample 20. With rising temperature, the combination of decreasing glass viscosity and increasing in-depth SiC oxidation can lead to greater gas formation rates below the glassy layer and, potentially, bubble formation. This interpretation is consistent with the surface SEM images of Figs. 6; the larger areas of exposed zirconia grains on sample 9 have an appearance that suggests the bursting of bubbles formed by gases evolving from beneath a relatively thick glassy layer, whereas on sample 20 the exposed areas are more consistent with direct oxidation of ZrB2 to ZrO2 on the outer surface. Comparison of furnace-oxidized specimens with Plasmatronexposed specimens suggests that the formation of silica is faster and/or the volatilization of silica is slower in the dissociated oxygen environment of the Plasmatron as compared with the furnace environment. Simple Gibb’s free energy calculations based on thermodynamic data from [47] show that the reaction of SiC with atomic oxygen has a larger thermodynamic driving force than the reaction with molecular oxygen. Similar thermodynamic computations show that SiO2 is higher in an atmosphere of 2 \\x02 103 Pa O2 (furnace environthe equilibrium partial pressure of SiO gas over than in an atmosphere of 3 \\x02 103 Pa O atoms (Plasmatron ment) environment) for all temperatures in our experiments. Experiments at \\x18900\\x0eC using a microwave-discharge source to generate atomic oxygen in a \\x18400 Pa O2 =Ar mixture have shown dramatically increased silica formation on Si wafers, and SiC and thin films, over similar experiments without atomic oxygen [48]. The O-atom densities in these discharge experiments were 2-4 orders of magnitude lower than in the Plasmatron tests. Balat et al. [49,50] have studied the active-passive oxidation transition boundary of silicon carbide in standard and microwaveexcited air. The theoretical transition temperature between active and passive SiC oxidation, computed using Wagner’s model for boundary-layer diffusion-limited surface oxidation [51], was determined to be the same for molecular and atomic oxygen when pressure (i.e., PO2 \\x88 0:5PO ) the latter was expressed as equivalent molecular oxygen partial [50]. Experimentally, however, they found that the temperature-pressure domain characterized by passive silica formation on sintered SiC specimens was significantly  Si3N3  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970   \\x0c', '276  MARSCHALL ET AL.  11 \\x16m),  the ZrB2  of ZrB2  ZrB2  -10HfB2  -  -15SiC and ZrB2  16-21 kJ \\x01 g\\x001 ,  enlarged over the theoretical domain by oxygen dissociation [49,50]. Both our highest-temperature furnace oxidation (at 1575\\x0eC) and the two highest-power Plasmatron tests (15-9 and 16-4) lie just inside the active SiC oxidation regime as calculated by Balat et al. [49,50]. Our finding of a glassy outer layer on sample 9 (Fig. 6) versus a surface on the 1575\\x0eC furnace-tested specimen largely zirconia (Fig. 10) is consistent with the hypothesis of an expanded passive oxidation domain under dissociated oxygen. In general, surface properties like emittance and catalytic efficiency depend on both the chemical nature and the microstructure of the surface. Based on spectral reflectance measurements, and the temperature-dependent total emittance behavior estimated from these data, oxidation in the Plasmatron increases the emittance of -30SiC materials. This result is in good agreement with the measurements of Scatteia et al. [32], who showed that the emittance -15SiC-2MoSi2 composites heated in an oxidizing environment was always higher than that of specimens heated under vacuum. Despite differences in the morphology of the posttest surfaces (e.g., Fig. 6), reflectance spectra varied little over most of the spectral range (except in the immediate vicinity of the Si-O stretching mode feature between 7 and leading to an essentially identical estimated temperature-dependant total emittance behavior for all posttest specimens. In a very recent study, Scatteia et al. [52] investigated the effects of diamond tool versus electrical discharge machining methods (producing different surface roughness) on the emittance of ZrB2 15SiC materials. They found little difference between the samples once the surfaces were oxidized, concluding that the chemical composition of the oxide was the dominant factor over surface morphology in determining emittance. Ito et al. [53] tested ZrB2 -based UHTC specimens in a flat-face stagnation point configuration similar to ours (see Fig. 1b) in a 110 kW ICP wind tunnel at gas enthalpies of reaching surface temperatures of between about 1400 and 1850\\x0eC. The samples became whitish after testing at their highest heating conditions. Ito et al. interpreted an apparent jump in sample surface enthalpy (near 19 kJ \\x01 g\\x001 ) temperature with increasing gas as a reflection of decreasing emittance and increasing catalytic efficiency due to sample oxidation. No sample manufacturing information or posttest characterization results were presented, though the whitish surface is identified by the authors as “oxidized zirconium,” suggesting that no silica former was incorporated. By contrast, the formation of a borosilicate glassy layer on the specimens tested at VKI resulted in a black posttest surface with relatively high emittance and moderate catalytic efficiency, as reflected by the lower steady-state surface temperatures for higher stream enthalpies. Monteverde and Savino [33] have recently tested a hot-pressed -15SiC composite at atmospheric pressure in an 80 kW plasma torch facility located at the University of Naples. Samples were radius. Tests were run at total enthalpies between 14 and 20 kJ \\x01 g\\x001 , diamond machined into hemispherical test specimens with a 7.5 mm reaching surface temperatures of up to 1930\\x0eC for up to 4.5 min. Posttest sample surfaces were smooth and dark in appearance. Sample emittance during testing, derived by simultaneous twoand one-color radiometry measurements, was about 0.9. Simulations of fully catalytic versus noncatalytic heating of the test specimens suggest that the oxidized UHTC surfaces must possess very low catalytic efficiency to explain the measured surface temperature transients. Both the optical and catalytic behaviors reported by Monteverde and Savino are generally consistent with those found in the present study. However, cross-sectional SEM images of their oxidized specimens reveal a somewhat different oxide structure, with the outer glassy layer and the virgin UHTC material separated by a distinct layer imbedded in glass over a thinner SiC-depleted layer. Similar multilayer oxide structures have been reported in furnace experiments and arcjet tests [14]. In the present study, only a single sublayer of partially oxidized UHTC is observed, perhaps because the temperatures were not sufficiently high and/or the exposure times were not sufficiently long to fully deplete the SiC from the layer beneath the outer glassy scale. In furnace  ZrB2  -30SiC  of ZrO2  oxidation tests, SiC depletion is only significant at test temperatures approaching 1500\\x0eC, but it occurs readily at higher temperatures [12]. The depletion of SiC from the layer beneath the outer glassy layer occurs by active oxidation of SiC to SiO and CO, which has a strong temperature dependence [15].  VI.  Conclusions  ZrB2  122:5 W \\x01 cm\\x002  Ultrahigh-temperature ceramics are candidates for leading-edge and control surface applications on future hypersonic aerospace vehicles, for which sustained operation at surface temperatures of 1500-2000\\x0eC and above are desired [1,2]. In such flight environments, materials will be exposed to highly dissociated air, particularly atomic oxygen. This study tested materials in a partially dissociated air environment at the lower end of the desired temperature range, but with generally favorable results. Under the conditions of this study, exposure to cold-wall heat fluxes ranging from 51.0 to resulted in surface temperatures ranging from 1240 to 1570\\x0eC. The Plasmatron stream consisted of highly dissociated oxygen and, as heat flux increased, an increasing proportion of dissociated nitrogen. The exposures resulted in oxidation of the specimen surfaces, which produced an outer layer of silica-based glass and an underlying layer of ZrO2 with some residual and SiC. Changes in specimen dimensions and mass were minimal. The derived catalytic efficiency of the oxidizing surface low (\\x1810\\x003 ) and the measured posttest emittance was high was (\\x180:9), both favorable properties that help lower the quasi-steady surface temperatures attained during hypersonic flight. Future experiments are planned using a redesigned model configuration, which will enable higher-enthalpy/higher-temperature studies. In addition, other variables, such as the volume fraction of SiC in the UHTC composite and the composition of the UHTC material (HfB2 vs ZrB2 ), will be evaluated to determine their effects on the properties of the resulting surfaces. Additional diagnostics are planned for these experiments, in particular a wideband radiometer to enable in situ determination of the surface emittance and emission spectroscopic diagnostics for the detection of volatile species. The importance of the former diagnostic is evident, given the significant uncertainty that emittance introduces into evaluations of the surface test environment and such properties as catalytic efficiency. The value of the latter diagnostic was recently demonstrated in an adjunct to this test series in which the temporal emission signatures of volatile boron species were detected during UHTC oxidation in the Plasmatron; these results will be reported in a separate paper [54].  Acknowledgments  This research was supported by the Ceramics and Nonmetallic Materials Program of the U.S. Air Force Office of Scientific Research through contracts F49550-05-C-0020 (Marschall and Pejaković) and FA9550-06-0125 (Fahrenholtz and Hilmas), the National Science Foundation through grants DMR-0435856 (Marschall) and DMR0346800 (Fahrenholtz and Zhu), and the European Office of Aerospace Research and Development through contract FA865506-1-3078 (Fletcher, Asma, and Thömel). 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[54] Playez, M., Fletcher, D. G., Marschall, J., Fahrenholtz, W. G., Hilmas, G. E., and Zhu, S., “Optical Emission Spectroscopy during Plasmatron Testing of ZrB2 -SiC Ultra-High Temperature Ceramic Composites,”  the American  and ZrB2  Journal  Journal of Thermophysics and Heat Transfer, Vol. 23, No. 2, 2009,  ZrB2  [53]  of  Polycrystalline SiC and LPCVD Si3N4 ,” High Temperature Corrosion  pp. 279-285.  Downloaded by Oregon State University on January 12, 2015 | http://arc.aiaa.org | DOI: 10.2514/1.39970 \\x0c']"
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  "_id": 185,
  "PDF": "Oxidation of ZrC–30 vol_ SiC composite in air from low to ultrahigh temperature.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 32 (2012) 947-954  Oxidation of ZrC-30 vol% SiC composite in air from low to ultrahigh temperature  Liyou Zhao, Dechang Jia  , Xiaoming Duan, Zhihua Yang, Yu Zhou  ∗  Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin 150080, China  Received 3 May 2011; received in revised form 13 October 2011; accepted 17 October 2011  Available online 8 November 2011  Abstract  Oxidation of ZrC-30 vol.% SiC is investigated in air using furnace and oxyacetylene torch. The microstructure and phase composition of oxide scales are analyzed via SEM, XRD, and Raman. At 800 and 1100 C, SiC is embedded in the porous and cracked ZrO2 scales, which have a single-layer structure and are almost non-protective. At 1300 and 1500 C, the protective effect of oxide scales is enhanced by the formed SiO2 . C for ≥1 h, and 1500 C for ≥15 min. The growth kinetics The scales consist of two subscales, outer and inner layers, during oxidation at 1300 of both layers is analyzed. At 1700 C, a new layer is observed between the outer and inner layers, which should contain less carbon than the inner layer. At 2100 C, the oxide scale is porous and contains many big holes. This scale shows a single-layer structure, which mainly consists of ZrO2 . © 2011 Elsevier Ltd. All rights reserved.                    Keywords: A. Hot pressing; D. Carbides; D. SiC; D. Carbon; Oxidation  1.  Introduction  Interest in ultrahigh-temperature ceramics (UHTCs) has increased significantly in recent years due to the drive to produce a thermal protection system and other components for hypersonic aerospace vehicles.1-3 Zirconium carbide (ZrC) is an important member of UHTCs. Besides its high melting temperature, ZrC has a unique combination of high fracture strength, high electrical and thermal conductivity, and resistance to erosion/corrosion.4-6 Oxidation resistance is a major issue in the development of UHTCs, but ZrC has a poor high-temperature chemical stability in oxidizing atmosphere, which significantly limits its actual application as UHTCs. Previous reports indicated that ZrO2 scales on ZrC can be divided into two layers.7-12 The outer layer is porous and cracked, containing a small amount of free carbon. Pores and cracks in the outer layer offer channels for inward diffusion of oxygen. So, the outer layer is non-protective. The inner layer is relatively dense and rich in carbon, which attracts  ∗  Corresponding author. Tel.: +86 451 86418792; fax: +86 451 86414291.  E-mail addresses: dcjia@hit.edu.cn, dechangjia@yahoo.com (D. Jia).  0955-2219/$ - see front matter © 2011 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2011.10.024  lots of attention in the past several years. This layer is considered as a barrier for the diffusion of oxygen during oxidation. However, the inner layer is very thin; that is to say, oxide scale on ZrC mainly consists of the porous and cracked outer layer. Cracks in the outer layer penetrate easily into the inner layer, accelerating the oxidation of matrix. So, the oxide scale on ZrC is almost non-protective, which is confirmed by the linear or approximately linear oxidation kinetics.7,9,13,14 A common approach to improve oxidation resistance of UHTCs is the incorporation of Si-containing compounds into UHTCs matrix to form a protective SiO2 -containing oxide scale.15-18 However, the reports are scarce by far on oxidation of ZrC-based composites with additives of Si-containing compounds, and some results deviate from that expected. Pierrat et al.19 investigated the oxidation resistance of pressurelesssintered ZrC-20 vol.% MoSi2 using an experimental facility called REHPTS. The results showed that addition of 20 vol.% MoSi2 was detrimental to the oxidation behavior of ZrC in the temperature range of 1800-2400 K, because of its dissociation and its role in the surface melting. Li et al.20 prepared a ZrC-SiC coating on 2D C/ZrC-SiC composites by chemical vapor deposition and investigated the oxidation behavior of this composites at 1800 C using CH4 combustion wind tunnel. The results     \\x0c', '948  L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954     indicated that no protective oxide scale was formed, because ZrC was oxidized completely very fast, and SiC or its oxide SiO2 was easily blown off. Zhao et al.21 prepared C/ZrC-SiC by polymer infiltration and pyrolysis process. Furnace test at 1200 C in air indicated that introduction of 1.9 vol.% SiC can improve the oxidation resistance of C/ZrC, due to formation of dense SiO2 glass. In this study, we investigated the oxidation behavior of hot-pressed ZrC-30 vol.% SiC in air from low to ultrahigh temperature using furnace and oxyacetylene torch. The structure evolution of oxide scales was described and discussed. Some interesting phenomena were reported which have not been found in previous works.  2. Material and methods        ZrC-30 vol.% SiC composite was referred as ZrC-30SiC below. The starting powders were ZrC (mean particle size 2.1 \\u242em, >98% purity, Changsha Wing High High-Tech New Materials Co., Ltd., China), SiC (␤-SiC, mean particle size 2 \\u242em, >99% purity, Central Iron & Steel Research Institute, China). The powder mixture of ZrC-30 vol.% SiC was ball milled in ethanol for 24 h and then dried. Mixed powders were then uniaxially hot pressed in boron nitride-coated graphite die at 2000 C for 60 min under an argon atmosphere with an applied pressure of 30 MPa. Bulk density and theoretical density were evaluated using the Archimedes method and the 4.9 mm × 5.8 mm × 6.6 mm were cut from the hot-pressed specrule of mixtures, respectively. Sample coupons in the size of imens, and all surfaces were diamond polished to a 1 \\u242em finish. Coupons were ultrasonically cleaned in acetone and alcohol, and then exposed to air at temperature of 800, 1100, 1300 and 1500 C, respectively, for 15 min to 4 h. The oxidation experiment was conducted in a box furnace with MoSi2 heating elements. The weight of samples before and after oxidation was carefully measured using a precision balance with an accuracy −5 g. Sample coupons with dimension of Φ17.4 × 12.7 mm of 10 were used in the oxidation experiment at temperature of 1700 and 2100 C, which was carried out with the oxyacetylene torch facility of Hu et al.’s group.22,23 The oxyacetylene torch test was conducted under two different conditions. During the first test, after several adjustments of gas flow rate, the pressure −1 , and and flux of acetylene were fixed at 0.1 MPa and 0.32 m3 h −1 , respectively. The test time for oxygen 0.5 MPa and 0.64 m3 h was about 13 min. The sample surface achieves a temperature of 1670 C within about 3 min, and then it gradually increases to 1700 C. During the second test, the pressure and flux of acety−1 , and for oxygen 0.5 MPa lene were 0.1 MPa and 0.75 m3 h −1 , and 1.5 m3 h respectively. The test time was about 10 min. The temperature of sample surface increases sharply to 2060 C and then gradually reaches to the maximal value of 2200 C. The phase composition of oxide scale was identified using Xray diffractometer and Raman spectroscopy. X-ray diffraction (XRD) was carried out in a D/max-2200VPC diffractometer (Rigaku, Tokyo, Japan) with Cu K␣ radiation. The Raman spectra were recorded using the 458 nm line from an argon ion laser using a Raman system (JY HR800, Paris, France). The surface               Fig. 1. SEM micrograph of the polished surface of ZrC-30SiC.  and cross section of oxidized samples were observed by scanning electron microscope (SEM, FEI Quanta 200F, Eindhoven, the Netherlands). An optical microscope (Axiovert 40MAT, Germany) was also employed to observe the cross section of oxide scale and determine its thickness.  3. Results  3.1. Density and microstructure  The measured bulk density of ZrC-30SiC is 5.56 g/cm3 , which corresponds to a relative density of 98.4%. Fig. 1 shows a SEM micrograph of the polished surface of ZrC-30SiC. The dark and grey phases are SiC, and they appear to be uniformly dispersed in the light ZrC matrix. The different contrast of SiC should be related to their different crystallographic orientations.24,25 Microstructure of the composite is regular, and few pores are observed on the polished surface, which supports the result of density measurement.  3.2. Oxidation at low temperature range (800 and 1100     C)             XRD analysis (Fig. 2a) indicates that the oxide scale formed at 1100 C for 15 min consists of m-ZrO2 . Fig. 3a shows the surface SEM micrograph of this scale. It is cracked and porous. SiC is embedded in ZrO2 . Because the strongest peak of SiC (1 1 1) at 35.6 and the peaks of m-ZrO2 (2 0 0) at 35.3 , ( ¯1 0 2) at 35.9 are too close to be distinguished, SiC is difficult to detect by XRD in m-ZrO2 matrix. Cross-section SEM micrograph (Fig. 3b) indicates that the oxide scale has a single-layer structure. Its thickness is about 290 \\u242em, and mass gain of sample is 177.9 g/m2 . Fig. 4 shows the Raman spectrum of this scale. −1 and 1600 cm −1 , corTwo broad peaks appear around 1350 cm responding to the A1g mode associated with amorphous carbon and E2g with graphite, respectively.8 Fig. 5 shows the specific mass-change as a function of exposure time for ZrC-30SiC during oxidation at 800 and 1100 C. The kinetics follows an approximately linear law, indicating that the oxidation proceeds mainly via interface limited reaction.     \\x0c', 'L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  949  m-ZrO2  SiO2  SiC  (f)  (e)  (d)  (c)  (b)  (a)  20  25  30  2 Theta (degree)  35  40  300  150  0  )  u  .  a  (  y  t  i  s  n e  t  n  I  1200  1500  1800  Raman shift (cm-1)  Fig. 4. Raman spectrum of the oxide scale on ZrC-30SiC oxidized at 1100     C        C for 15 min,  for 15 min.     Fig. 2. XRD patterns of the oxide scale on ZrC-30SiC: (a) 1100     (b) 1300 C for 15 min, (c) outer layer at 1500 C for 15 min, (d) inner layer at C for 15 min, (e) the first oxyacetylene torch test (1700 1500 second oxyacetylene torch test (2100 C).  C) and (f) the        3.3. Oxidation at intermediate temperature range (1300 and 1500 C)              Microstructure and phase composition of the oxide scale formed at 1300 C for 15 min are similar to that of scale formed at 1100 C for 15 min. No SiO2 is detected on the surface of oxidized sample by XRD (Fig. 2b) and SEM (Fig. 6a). The oxide scale shows a single-layer structure (Fig. 6b). The thickness of this scale is about 295 \\u242em; mass gain of sample is 179.8 g/m2 . After oxidation at 1500 C for 15 min, most SiC on the surface of sample oxidizes to form SiO2 (Fig. 6c). The oxide scale shows a duplex structure (Fig. 6d). The thickness of this scale is about 300 \\u242em, and mass gain of sample is 181.6 g/m2 . High magnification micrographs indicate that, the outer layer contains some pores, and the grey SiO2 mainly distributes in these pores; the inner layer appears to be dense, and SiC is almost not oxidized in the inner layer. For further analyzing the phase composition, each layer of scale was separated from sample. The outer layer is white in color. XRD pattern (Fig. 2c) reveals  0.9  0.6  ) 2  m  /  g  k  (  m  0.3  0.0  1100 °C  800 °C  0  80  160  240  ( min)  Fig. 5. Specific mass-change  (\\x01m)  as  a  function of  exposure  time  (t)  for  ZrC-30SiC during oxidation at 800 and 1100  C.     that it consists of m-ZrO2 and SiO2 . Raman spectrum analysis, shown in Fig. 7a, indicates presence of some free carbon in the outer layer. The Raman shifts between 100 and 1000 cm are characteristic peaks for m-ZrO2 and possibly SiO2 .26,27 The inner layer is black in color. XRD pattern (Fig. 2d) reveals that  −1  Fig. 3. Surface (a) and cross-section (b) SEM micrographs of ZrC-30SiC oxidized at 1100     C for 15 min.      \\x0c', '950  L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  Fig. 6. Surface and cross-section SEM micrographs of ZrC-30SiC oxidized at 1300 (a and b) and 1500     C (c and d) for 15 min.  the matrix of inner layer is composed of m-ZrO2 . Raman spectrum of inner layer (Fig. 7b) shows much bigger intensity ratio of the peaks for free carbon to the peaks for m-ZrO2 than the spectrum of outer layer, indicating that the inner layer contains more free carbon. Similar duplex-structure scales are found durC for ≥1 h, and 1500 ing oxidation of ZrC-30SiC at 1300 C for ≥15 min.                 The time dependency of thickness of the inner and outer layers during oxidation at 1300 and 1500 C is shown in Fig. 8. At both temperatures, the thickness of outer layer almost does not change with time after forming duplex structure, while the thickening of inner layer is approximately parabolic with time. The outer layer is thicker at 1300 C than at 1500 C. Fig. 9 shows the specific mass-change and its corresponding square as a function of exposure time for ZrC-30SiC during oxidation at 1300 and 1500 C. The kinetics is in approximate agreement with parabolic law, indicating that the oxidation is mainly controlled by diffusion. Fig. 10 shows the optical macrograph of ZrC and ZrC-30SiC oxidized at 1300 and 1500 C for 15 min. We can see that the ZrC samples are oxidized catastrophically while the oxidation of ZrC-30SiC is not so serious. ZrC-30SiC has a better oxidation resistance than ZrC in air at 1300 and 1500 C.           3.4. Oxidation at high and ultrahigh temperature (about 1700 and 2100 C)     Fig. 11a shows a surface SEM micrograph of the scale on sample after the first oxyacetylene torch test (1700 C). Many pores and a small amount of SiO2 are observed. SiO2 is too little to be detected by XRD (Fig. 2e). Low magnification SEM micrograph (Fig. 11b) indicates that the cross-section morphology of this scale is similar to that of scale formed at 1500 C for 15 min, except formation of an interlayer between the outer        Fig. 7. Raman spectra for the outer (a) and inner (b) layers of oxide scale on  ZrC-30SiC oxidized at 1100  C for 15 min.     \\x0c', 'L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  951  Inner layer Outer layer  1300 °C  )  m  µ  (  d  )  m  µ  (  d  450  300  150  0  1200  800  400  0  18  12  6  0  0  4  0 1  1  102min  2  3  80  t (min)  160  240  1500 °C  3  Inner layer  Outer layer  2  1  102min  0  4  0 1  9  6  3  0  0  Fig. 10. Optical macrograph of ZrC and ZrC-30SiC oxidized at 1300 and     1500  C for 15 min.     mainly composed of ZrO2 and SiC; the amount of SiO2 in the outer layer is decreased, and porosity increased, compared with that of the outer layer during oxidation at 1500 C for 15 min. Based on the analysis in Section 3.3, it can be inferred that the carbon content of interlayer is lower than that of inner layer, but higher than that of outer layer. After the second oxyacetylene torch test (2100 C), many macroscopic holes are present on the surface of scale, and this scale mainly consists of m-ZrO2 (Fig. 2f). Surface SEM micrograph (Fig. 11d) shows three different regions, namely, porous ZrO2 skeleton, big holes, and a little SiO2 . This scale has a single-layer structure (Fig. 11e), and its thickness is about 350 \\u242em.     4. Discussion  We know that one of the differences between the two layers of oxide scales on ZrC-30SiC is the higher carbon content of inner layer, which is consistent with the previous reports on oxidation of ZrC. During XRD analysis, no obvious ZrO2 peak shift is observed in the pattern of inner layer when compared with that of outer layer, but we do not exclude the possibility of presence of reported oxycarbide (ZrO2−xCy here) in the inner layer.28,29 Carbon is easy to burn out in air at high temperature, whether in the lattice of ZrO2−xCy or as a simple substance.30-32 So, it depends on the protective effect of oxide scales whether the inner layer appears. Oxidation rate of SiC is negligible below 1100 C in air.30 Oxidation behavior of ZrC-30SiC should be similar to that of ZrC at this temperature range. Generally, dense ZrC C in air,7 and the oxidation ceramic oxidizes remarkably at 800 by reaction (1) is thermodynamically the most probable.33 ZrC(s) + O2 (g) = ZrO2 (s) + C(s)  (1)        The porous and cracked oxide scales are almost nonprotective during oxidation at 800 and 1100 C. Most of the formed carbon would be oxidized to generate gases CO2 and CO. The oxide scales show a single layer structure. Because     0  80  160  t (min)  240  Fig. 8. Time (t) dependency of the thickness (d) of inner and outer layers during is the plot of d2 vs t for the data of  oxidation at 1300 and 1500  C. The inset     inner layer.  and inner layers. Porosity and color (Fig. 11c) of the interlayer are intermediate between that of the outer and inner layers. The thickness of oxide scale is about 220 \\u242em. High magnification micrographs indicate that, both inner layer and interlayer are  Fig. 9. Specific mass-change (\\x01m) and its corresponding square (inset) as a  function of exposure time (t)  for ZrC-30SiC during oxidation at 1300 and     1500  C.      \\x0c', '952  L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  Fig. 11. Surface and cross-section SEM micrographs of ZrC-30SiC after the first (1700 Shows a corresponding optical micrograph of (b).     C, a and b) and second (2100     C, d and e) oxyacetylene torch test. (c)        diffusion of CO2 and CO through ZrO2 is difficult,9 a small amount of carbon remains in the scale. SiC starts to oxidize significantly at 1300 C.34 The formed SiO2 could seal the pores and cracks of scales, increasing the density of scales thus retarding inward diffusion of oxygen. However, the oxide scale formed at 1300 C for 15 min still can not protect most of the formed carbon from oxidation, because the exposure time is too short to generate enough SiO2 . In contrast, during oxidation at 1500 C for 15 min, more SiO2 is formed due to the more rapid oxidation of SiC. The local oxygen concentration in the regions of the matrix/scale interface becomes so low that a carbon-rich ZrO2 layer is formed, i.e., the observed black inner layer of scale. Growth process of each layer of scales on ZrC-30SiC during oxidation at 1300 and 1500 C can be described as follows. ZrC oxidizes quickly to form a porous ZrO2 scale, i.e., the outer layer. The thickness of outer layer increases with time. SiC oxidizes relatively slowly to form SiO2 and enhances the protective effect of outer layer. As the inner layer forms, the outer layer stops thickening and the inner layer thickens with time. Oxidation rate of SiC is relatively lower at 1300 C. So, the inner layer appears later, and the outer layer is thicker. As mentioned above, carbon is chemically unstable in air at high temperature. The outer layer of scales is relatively porous. However, why doesn’t the carbon of inner layer near the interface of two layers oxidize gradually to result in thickening of outer layer? A           rational conjecture is that, the formed carbon diffuses outward and reacts with oxygen in the outer layer; also, oxidation of carbon is slower than diffusion. Outward diffusion of carbon has been observed during oxidation of ZrC and other carbides.7,35,36 The outer layer of scales on ZrC-30SiC is protective, which can lower the oxidation rate of carbon. During oxidation, lots of inward diffused oxygen is consumed in the outer layer by reaction with outward diffused carbon. Due to the sufficient carbon from inner layer, the interface between inner and outer layers can not move inward. Zr is thermodynamically oxidized more easily than carbon.33 Some oxygen diffuses through the inner layer to react with ZrC-30SiC matrix, resulting in thickening of the inner layer. Previous reports indicated that the porous outer layer thickens linearly with time, while the dense inner layer parabolically and then attains a limited constant thickness during oxidation of ZrC and other refractory carbides at low oxygen pressure.9,37 The amount of carbon created at ZrC matrix/scale interface decreases with increasing of oxygen pressure.38 So, the inner layer of oxide scale is usually difficult to observe during oxidation of ZrC in air.7,14,39 In this study, ZrC forms a nonadherent oxide scale below 1500 C which is brittle to touch, as Fig. 10 shows, and we do not find obvious duplex structure in them. Introduction of SiC enhances the protective effect of outer layer of scales on ZrC, promotes formation of the dense inner layer, and increases the volume ratio of inner layer to outer     \\x0c', 'L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  953  layer, therefore improving the oxidation resistance of ZrC in air at intermediate temperature range. SiC starts to oxidize actively at 1650 (2).40  C according to reaction     SiC(s) + O2 (g) = SiO(g) + CO(g)  (2)     The amount of SiO2 in the scale on ZrC-30SiC would decrease with increasing of oxidation temperature.23,34,41 Dur(1700 ing the first oxyacetylene torch test C), lots of gas products can be generated due to the serious oxidation of ZrC and active oxidation of SiC, which leave behind many pores after passing through oxide scale. The protective effect of outer layer of oxide scale will be weakened. So, we assume that some carbon of inner layer in the regions near outer layer is oxidized, resulting in formation of a new layer with higher porosity and lighter color than the inner layer. During the second oxyacetylene torch test (2100 C), more gas products can be generated, which accelerates the damage of oxide scale, resulting in oxidation of the most formed carbon. The temperature under oxide scale is believed to be lower than the surface temperature. Part of SiC oxidizes passively to form SiO2 , which is beneficial to enhancing the oxidation resistance of ZrC.     5. Conclusions              Oxidation of ZrC-30 vol% SiC is studied in air. At 800 and 1100 C, the porous and cracked scales are almost nonprotective, and exhibit a single-layer structure. These scales consist of ZrO2 , SiC, and a small mount of carbon. At 1300 and 1500 C, SiO2 enhances the protective effect of ZrO2 scales. The ZrO2 C for ≥1 h, and 1500 scales show a duplex structure during oxidation at C for ≥15 min. The outer layer is 1300 white and relatively porous, containing SiO2 and less carbon; the inner layer is black and dense, containing SiC and more carbon. After forming duplex structure, the thickness of outer layer does not change with time, while the thickening of inner layer is approximately parabolic with time. At 1700 C, an interlayer is observed between the outer and inner layers of scale, which should contain less carbon than the inner layer. At 2100 C, the oxide scale is porous and contains many big holes. This scale shows a single-layer structure, and mainly consists of ZrO2 .        Acknowledgements  This work was supported by the Program for Changjiang Scholars and by the National Natural Science Foundation of China under grant No. 51021002.  3. Zou LH, Wali N, Yang J-M, Bansal NP. Microstructural development of a  Cf /ZrC composite manufactured by reactive melt infiltration. J Eur Ceram Soc 2010;30:1527-35.  4. Pierson HO. Handbook of  refractory carbides and nitrides: properties,  characteristics, processing, and application. Westwood, NJ: Noyes publi cations; 1996.  5. 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Properties of ceramics in the system ZrB2 -Ta5 Si3 . J Mater Res 2006;21:2593-9. 19. Pierrat B, Balat-Pichelin M, Silvestroni L, Sciti D. High temperature oxida tion of ZrC-20% MoSi2 in air for future solar receivers. Sol Energy Mater Sol Cells 2011;95:2228-37.  20. Li HB, Zhang LT, Cheng LF, Wang YG. Oxidation  analysis  of  2D  C/ZrC-SiC composites with different coating structures in CH4 combustion gas environment. Ceram Int 2009;35:2277-82.  21. Zhao D, Zhang CR, Hu HF, Zhang YD, Wang QK, He XB. Effect of SiC  on the oxidation resistance performance of C/ZrC composite prepared by polymer infiltration and pyrolysis process. Mater Sci Forum 2011;675677:415-8.  22. Zhou SB, Li WJ, Hu P, Hong CQ, Weng L. Ablation  behavior  of  ZrB2 -SiC-ZrO2 ceramic composites by means of the oxyacetylene torch. Corros Sci 2009;51:2071-9.  23. Han  JC, Hu P, Zhang XH, Meng SH, Han WB. Oxidation-resistant C. Compos Sci Technol 2008;68:799-806.  ZrB2 -SiC composites at 2200 24. Sciti D, Brach M, Bellosi A. Long-term oxidation behavior and mechani    cal strength degradation of a pressurelessly sintered ZrB2 -MoSi2 ceramic. Scripta Mater 2005;53:1297-302.  25. Ni D-W, Zhang G-J, Xu F-F, Guo W-M. Initial stage of oxidation process     C.  References  and microstructure analysis of HfB2 -20 vol.% SiC composite at 1500 Scripta Mater 2011;64:617-20.     1. Opeka MM, Talmy IG, Zaykoski  JA. Oxidation-based materials  selec tion for 2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci 2004;39:5887-904.  2. Levine SR, Opila EJ, Halbig MC, Kiser JD, Singh M, Salem JA. Evaluation  of ultra-high temperature ceramics for aeropropulsion use. J Eur Ceram Soc 2002;22:2757-67.  26. Djurado E, Bouvier P, Lucazeau G. Crystallite size effect on the tetragonal monoclinic transition of undoped nanocrystalline zirconia studied by XRD and Raman spectrometry. J Solid State Chem 2000;149:399-407.  27. Berry FJ, Skinner SJ, Bell IM, Clark RJH, Ponton CB. The influence of pH  on zirconia formed from zirconium (IV) acetate solution: characterization  by X-ray powder diffraction and Raman spectroscopy. J Solid State Chem 1999;145:394-400.  \\x0c', '954  L. Zhao et al. / Journal of the European Ceramic Society 32 (2012) 947-954  28. Bargeron CB, Benson RC. X-ray microanalysis of a hafnium carbide film oxidized at high temperature. Surf Coat Technol 1988;36:111-5.  29. Bargeron CB, Benson RC, Jette AN, Phillips TE. Oxidation of hafnium     C.  J  Am Ceram Soc  carbide in the 1993;76:1040-6.  temperature  range  1400-2060  30. Fahrenholtz WG.  Thermodynamic  analysis  of  ZrB2 -SiC  formation  of  a  SiC-depleted  region.  J  Am  Ceram  Soc  oxidation: 2007;90:  143-8.  35. Zhang JF, Wang LJ, Jiang W, Chen LD. High temperature oxidation behav ior and mechanism of Ti3 SiC2 -SiC nanocomposites in air. Compos Sci Technol 2008;68:1531-8.  36. Reichle M, Kickl JJ. Investigation about the high temperature oxidation of titanium carbide. J Less-Common Met 1972;27:213-36.  37. Shimada S, Yunazar F. Oxidation of hafnium carbide and titanium carbide  single crystals with the formation of carbon at high temperatures and low oxygen pressures. J Am Ceram Soc 2000;83:721-8.  31. Wang DN, Liang KM. The effect of carbon on the phase stability of zirconia. J Mater Sci Lett 1998;17:343-4.  38. Rao GR, Venugopal V. Kinetics and mechanism of the oxidation of ZrC. J Alloys compd 1994;206:237-42.  32. Heuer AH, Lou VLK. Volatility diagrams for silica, silicon nitride, and  39. Sciti D, Silvestroni L, Guicciardi S, Fabbriche DD, Bellosi A. Processing,  silicon carbide and their application to high-temperature decomposition and oxidation. J Am Ceram Soc 1990;73:2789-803.  33. Shimada  S. A thermoanalytical  study  on  the  carbon. Solid State  oxidation of ZrC and Ionics 2002;149:  HfC powders with formation of  319-26.  34. Zhang X-H, Hu P, Han J-C. Structure evolution of ZrB2 -SiC during the oxidation in air. J Mater Res 2008;23:1961-72.  mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2 . J Mater Res 2009;24:2056-65. 40. Vaughn WL, Maahs HG. Active-to-passive transition in the oxidation of silicon carbide and silicon nitride in air. J Am Ceram Soc 1990;73:1540-3.  41. Han W-B, Hu P, Zhang X-H, Han J-C, Meng S-H. High-temperature oxida tion at 1900 C of ZrB2 -xSiC ultrahigh-temperature ceramic composites. J Am Ceram Soc 2008;91:3328-34.     \\x0c']"
},{
  "_id": 186,
  "PDF": "Oxidation resistance and microstructure evolution of ZrB2–SiC–La2O3-SiC dual-layer coating on siliconized graphite at 1800 °C under low air pressures.pdf",
  "Text": "['Ceramics International 46 (2020) 27150-27157  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Oxidation resistance and microstructure evolution of ZrB2-SiC-La2O3/SiC dual-layer coating on siliconized graphite at 1800 °C under low air pressures  T  Yan Rena,b, Yuhai Qiana,∗, Jingjun Xua, Jun Zuoa, Meishuan Lia  a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China b School of Materials Science and Engineering, University of Science and Technology of China, Shenyang, 110016, China  A R T I C L E  I N F O  A B S T R A C T  The oxidation behaviors of a ZrB2-SiC-La2O3/SiC dual-layer coating on siliconized graphite at 1800 °C under low air pressures (50, 5 and 0.5 kPa) were investigated. The results showed that with the decrease of air pressure, the oxidation kinetics of the coated samples changed from parabolic weight gain to linear weight loss. A protective oxide scale consisted of ZrO2 and SiO2 with La dispersed was formed on the coating surface after oxidation in 50 kPa air. The oxide scale formed in 5 kPa air was full of bubbles. Only porous ZrO2 layer was left on the coating surface after oxidation in 0.5 kPa air. At 1800 °C, the active oxidation of SiC occurred and gaseous SiO formed at the coating/oxide interface. The surface volatilization of SiO became severe with the decrease of air pressure, resulting in the presence of non-protective oxide scale.  Keywords: Coating Oxidation resistance Graphite Low pressure air  1.  Introduction  Carbon materials gain wide application in hypersonic flight vehicles because of their attractive high-temperature mechanical properties [1]. However, poor oxidation resistance at elevated temperatures in the oxidizing environment severely limits their applications [2]. Applying ceramic coatings is a practical strategy to improve the oxidation resistance of carbon materials [3,4]. Ultra-high temperature ceramics (UHTCs) with good oxidation resistance and high thermal stability are potential materials toward anti-oxidation coatings [5,6]. Among UHTC family, ZrB2 is extensively concerned due to its attractive comprehensive performance. Nevertheless, the protective performance of ZrB2 is limited up to 1200 °C due to the rapid volatilization of liquid B2O3, resulting in a significant decrease of oxidation resistance. The introduction of silicides could enhance the protection capability of ZrB2 significantly by forming molten SiO2 to seal or fill numerous open holes, the rapid diffusion channels of oxygen [7-10]. For these candidates of silicides, SiC is widely used as an effective additive toward diborides. Meanwhile, rare earth oxides, transition metals and transition metal silicides are also beneficial additives for ZrB2-SiC system to improve its oxidation resistance [11-13]. The hypersonic flight vehicles usually operate in near space with high Mach numbers. Under severe operating conditions, vehicles undergo ultra-high temperature and low-pressure atmosphere (correspondingly low oxygen partial pressure) during its flight and/or re entry process [14]. SiC oxidizes in an active mode at relatively high temperature and low oxygen partial pressure, leading to different oxidation behaviors of ZrB2-SiC system compared with that in atmospheric air. For example, Gao et al. [15] reported that when the ZrB2-SiC ceramic was oxidized in low-pressure O2/N2 mixture gas, the ZrSiO4 was promoted to form at below 1600 °C. Tian et al. [16] reported that during the oxidation of ZrB2-SiC composite in low-pressure O2/N2 mixture gas at 1500 °C, a transition of oxidation kinetics from parabolic law to linear law occurred with the decrease of oxygen partial pressure. This transition phenomenon was also found by Jin et al. [17] when ZrB2-SiC-Graphite was oxidized at 1800 °C in low-pressure O2/Ar mixture gases. Recently, Yang et al. [18] reported that when ZrB2-SiC-MoSi2 composite was oxidized in reduced pressure air at 1800 °C, the protective silica scale disappeared when air pressure was low enough. Therefore, active oxidation of SiC was more prone to occur at high temperatures and low oxygen partial pressures. Unfortunately, all of the above investigations were performed on UHTC bulks. To evaluate the service life of UHTC coating, the effect of air pressure on the oxidation behavior of UHTC coating is extremely essential to be understood. However, few related reports can be available. A ZrB2-SiC-La2O3/SiC dual-layer coating was prepared by a twostep method, i.e., slurry brushing and subsequent siliconizing in our previous work [19]. Under the condition at 1800 °C in 1 atm static air, the coating exhibited good protection capability by forming a continuous oxide scale. In this work, to further reveal the effect of air  ∗ Corresponding author. E-mail address: yhqian@imr.ac.cn (Y. Qian).  https://doi.org/10.1016/j.ceramint.2020.07.195 Received 12 May 2020; Received in revised form 19 July 2020; Accepted 19 July 2020  Available online 30 July 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', 'Y. Ren, et al.  Ceramics International 46 (2020) 27150-27157  Fig. 1. (a) Cross-sectional morphology of coated graphite sample, and XRD patterns of (b) the raw powders and (c) the outer layer of the as-prepared coating.  pressure on ultra-high temperature oxidation performances of such coating, oxidation tests were conducted at 1800 °C under low air pressures (50, 5 and 0.5 kPa), oxidation kinetics were determined, and microstructure evolution of the surface oxide scales was discovered. Finally, the oxidation mechanism and degradation process of the coating in the condition of ultra-high temperature and low air pressure were discussed.  then down to room temperature in about 10 min. In the whole oxidation period, air flowed evenly in the testing chamber, and constant pressure was maintained until the test was finished. To obtain oxidation kinetics (weight change per unit area vs. oxidation time), the dimensions and weight of the tested sample before and after oxidation test were measured by using a digital caliper and an electronic balance, respectively.  2. Experiment  2.1. Coating preparation  The square substrates with dimensions of 10 mm × 10 mm × 5 mm used in this study were cut from high-strength graphite. These graphite substrates were coated with SiC transition layers by siliconizing. The pack powders for siliconizing were composed of 62 wt% Si (300 mesh), 20 wt% C (300 mesh) and 18 wt% Al2O3 (300 mesh). The purpose of adding Al2O3 powders in the pack powders was to promote diffusion reaction and to avoid solidification of raw powders [20]. The siliconization was performed in a pressureless sintering furnace (graphite as heater) at 1600 °C for 2 h in Ar atmosphere. ZrB2-SiC-La2O3 The inner layer was prepared directly on SiC transition layers by slurry method. The slurry consisted of ZrB2 powder (38 wt%), SiC powder (6 wt%), La2O3 powder (6 wt%), polyvinyl butyral powder (1 wt%) and alcohol (100 wt%). After stirring magnetically for 30 min, they were brushed on SiC transition layers. After drying at 60 °C for 4 h, the SiC outer layer was prepared on the preliminary slurry layer by subsequent siliconizing, which was conducted at 1500 °C in argon atmosphere for 60 min. The pack powders were the same as those used to prepare the SiC transition layer.  2.2. Oxidation test  Oxidation test was performed in a high-frequency induction heating facility in our laboratory [21]. By using this facility, ultra-high temperature, extremely rapid heating and cooling, and controllable atmosphere could be achieved. Firstly, the sample was placed on a columnar graphite holder positioned in the center of an induction coil. Then the pressure of air in the chamber was adjusted and maintained by balancing the introducing air speed and exit pumping speed, which was monitored by a digital flowmeter. The pressures of air in the test were kept as 50, 5 and 0.5 kPa. Once the pressure of air was stable, the sample was heated up to 1800 °C within 3 min. The surface temperature of tested sample was monitored by Marathon MR1SCSF infrared twocolor pyrometer (Raytek, U.S.A.) with an accuracy of 0.75%. During the oxidation, the output power of electric source was adjusted frequently to keep desired surface temperature. After holding for required time, the tested sample was cooled down to 800 °C at a mean rate of ~10 °C/s  2.3. Characterizations of  the coating  The crystalline phases of the coating and oxidation products were determined by X-ray diffraction (XRD, Rigaku, Japan). The morphologies of the tested sample were observed by scanning electron microscope (SEM, LEO, Germany) equipped with energy dispersive spectroscopy analyzer (EDS, Oxford Instruments, UK). For cross-sectional analysis, the tested sample was mounted in resin and then ground and polished.  3. Results  3.1. Microstructure and phase compositions of  the coating  From the cross-sectional view of the coated sample (Fig. 1(a)), it can be recognized easily as four different regions: graphite substrate, SiC transition layer, and main coating with duplex structure (UHTC inner layer + SiC outer layer). The SiC transition layer was uneven, some residual C existed in it, and the graphite/SiC interface was undulated. The raw powders and outer layer of the main coating were analyzed by XRD, as shown in Fig. 1(b) and (c). Obviously, the inner layer was mainly composed of ZrB2, SiC and La2O3. Major SiC and minor Mullite (3Al2O3⋅2SiO2) were detected on the surface of the main coating. Inevitably, a trace of impurity O2 existed in the flowing Ar or was absorbed on the surface of pack powders. During the subsequent siliconizing process, the Al2O3 and Si in the pack powders reacted with impurity O2 to form minor Mullite. The Mullite in the SiC outer layer is beneficial to improve the oxidation resistance of the coating because it can react with SiO2 to form silicates glass with high viscosity and promote the formation of a crack-free oxide scale during oxidation [22,23]. The thicknesses of the inner and outer layers were about 50 μm, the total thickness of the main coating was about 100 μm.  3.2. Oxidation performance of  the coating  3.2.1. Oxidation kinetics Fig. 2 depicts the oxidation kinetics of coated graphite samples after oxidation at 1800 °C for 15 min in 50, 5 and 0.5 kPa air (the oxidation kinetics of coated samples after oxidation for 25 min was provided in Appendix A). The coated sample exhibited a slight weight gain of  27151  \\x0c', 'Y. Ren, et al.  Ceramics International 46 (2020) 27150-27157  Table 1  Detected phases by XRD on the sample surfaces after oxidation in air with different air pressures.  Air Pressure  50 kPa  5 kPa  0.5 kPa  Detected phases  SiC, Mullitem, aSiO2  ZrO2, SiC, Mullitem, a-SiO2  ZrO2, Mullitem, LaAl11O18  m  Note: m-minor.  3.2.3. Surface morphologies of the coating after oxidation Fig. 4 shows the macrographs of the surfaces of the coated samples after oxidation at 1800 °C for 15 min under different air pressures. The surface color of coated sample changed from completely dark to offwhite with the decrease of air pressure. Some bubbles were observed on the coating surface in 50 kPa air. Particularly, numerous white oxide whiskers formed on the coating surface in 5 kPa air, which were identified as amorphous SiO2 by EDS combined with XRD result. Moreover, spallation of part of the coating on the sample sides was observed in 0.5 kPa air. Fig. 5 presents the surface views of the coating after oxidation under different air pressures. The EDS analysis results of the local points marked in Fig. 5 were presented in Table 2. As shown in Fig. 5(a & b), the coating surface was mainly covered by continuous SiO2 glass after oxidation in 50 kPa air, and small bubbles as well as holes were also observed. As for 5 kPa air (Fig. 5(c & d)), numerous bubbles were observed on the coating surface, and a great number of small SiO2 wires existed on the top surface, stacked with a loose structure. Fig. 5(d) reveals that the oxide scale on the coating surface was not intact, some parts of it spalled. Based on the EDS results of points B and C, the surface oxide was mainly SiO2 glass, and the inner layer of the main coating was exposed and oxidized in the spallation regions of SiO2 glass. Particularly, some fine SiO2 wires distributed radially along the border of the exposed region, which suggested that the inner layer of the main coating was exposed because of the burst of bubbles. As for 0.5 kPa air (Fig. 5(e)), the coating surface was flat after oxidation. However, Fig. 5(f) shows clearly that the surface was covered with stacked white ZrO2 particles according to the EDS analysis result of local point D. Numerous holes with size of about 2-10 μm were observed, which provided rapid channels for inward diffusion of oxygen. Meanwhile, some parts of the coating surface exhibited different contrast (dark grey regions marked as dotted red lines). Combined with the EDS analysis result of local point E and XRD pattern (Fig. 3), the dark grey regions were mainly composed of Mullite and LaAl11O18.  3.2.4. Cross sections of the coating after oxidation Fig. 6 present the cross-sectional views of the coating after oxidation under different air pressures. For the case of 50 kPa air (Fig. 6(a & b)), the main coating still kept good bonding with SiC transition layer and penetrating cracks were absent. The SiC outer layer was oxidized and a continuous SiO2 scale formed, accompanied with some pores. Fig. 6(c) shows EDS element mappings of the marked rectangle Region 1 in Fig. 6(b), it can be seen that only part of the coating was oxidized, the SiC transition layer as well as graphite still kept intact. Meanwhile, La2O3 and Mullite existed as dispersed phases in SiO2 scale. On the contrary, the coating separated from SiC transition layer after oxidation in 5 kPa air (Fig. 6(d)). Enlarged views of two selected regions (marked as Region 2 and Region 3 in Fig. 6(d)), in Fig. 6(e & f), showed that the inner layer of the main coating was intact, but parts of the outer layer spalled, accordingly, the inner layer was exposed directly in these spallation zones. Furthermore, it can also be seen that a rugged SiO2 scale covered on the surface of the inner layer of the main coating, but there were many large holes in the remaining outer layer. On the top surface, fine wire-like SiO2 could be observed clearly, and they distributed loosely. As for 0.5 kPa air, the remaining coating and oxide scale detached from SiC transition layer, and a large gap (filled with  Fig. 2. Oxidation kinetics of different air pressures.  the coated samples at 1800 °C for 15 min under  1.15 × 10−3 g/cm2 after oxidation at 1800 °C for 15 min in 50 kPa air. According to the inset in Fig. 2, the oxidation gain curve in 50 kPa air was approximately parabolic one. On the contrary, the coated sample exhibited weight loss during the oxidation in 0.5 and 5 kPa air, and the final weight loss of the coated sample was 33.69 × 10−3 g/cm2 and 15.36 × 10−3 g/cm2, respectively. Notable weight loss suggested that a great number of gaseous products might be formed and volatilized during the oxidation process.  3.2.2. Phase compositions of oxide scales Fig. 3 presents XRD patterns of the surfaces of coated samples after oxidation at 1800 °C for 15 min under different air pressures (0.5, 5 and 50 kPa). The identified phases were listed in Table 1. The detected predominant phase on the sample oxidized in 50 kPa air was SiC, and minor Mullite was also detected. The existence of residual SiC indicated that the coating suffered from oxidation partly. Amorphous broad hump was observed, indicating that amorphous SiO2 was formed on the coating surface. In the case of 5 kPa air, besides SiC and Mullite, ZrO2 became predominant phase, and amorphous broad hump was also observed obviously. As for 0.5 kPa air, main ZrO2 and minor Mullite were identified, moreover, LaAl11O18 also appeared. Especially, no remnant SiC and amorphous SiO2 were detected under this condition, indicating that SiC was oxidized more severely, and silicon-containing oxides formed during oxidation and escaped into ambient surrounding as volatile gaseous species. Furthermore, no boron oxide was detected under all oxidation conditions, suggesting that boron oxide either dissolved into amorphous phase or volatilized into ambient surrounding immediately once it formed.  Fig. 3. XRD patterns of the surfaces of coated 1800 °C for 15 min under different air pressures.  samples  after  oxidation  at  27152  \\x0c', 'Y. Ren, et al.  Ceramics International 46 (2020) 27150-27157  Fig. 4. Macrographs of the surfaces of the coated samples after oxidation at 1800 °C for 15 min under different air pressures.  Fig. 5. Surface views of the coating after oxidation at 1800 °C for 15 min in (a, b) 50, (c, d) 5 and (e,  f) 0.5 kPa air.  resin) was left (Fig. 6(g)). Enlarged views of two selected regions (marked as Region 4 and Region 5 in Fig. 6(g)), in Fig. 6(h & i), showed that the SiC transition layer was also oxidized, and its surface became loose. The whole main coating was oxidized, and a porous ZrO2 layer  was left. Particularly, any Si-containing phase was not detected, indicating that Si was totally depleted due to the formation of Si-containing volatile gases during oxidation.  27153  \\x0c', 'Y. Ren, et al.  Table 2  EDS analysis results of  the marked points in Fig. 5.  O K (at%)  Al K (at%)  Si K (at%)  Zr L (at%)  La L (at%)  A B C D E  49.89 60.39 55.60 66.98 57.10  1.43 1.57 21.81 0.52 32.04  48.68 38.04 9.64 -  4.07  - -  12.95 32.42 4.43  - - -  0.08 2.36  4. Discussion  4.1. Volatility diagram for ZrB2-SiC at 1800 °C  For the tested coating, the following reactions might occur during oxidation in static air at 1800 °C.  2/3SiC(s) + O2(g) → 2/3SiO2(l) + 2/3CO(g)  2/5ZrB2(s) + O2(g) → 2/5ZrO2(s) + 2/5B2O3(l)  B2O3(l) → B2O3(g)  SiO2(l) → SiO2(g)  (1)  (2)  (3)  (4)  Reaction (1) is termed as ‘passive oxidation’ of SiC. However, SiC would be oxidized in an active regime to form volatile SiO(g) when the  Ceramics International 46 (2020) 27150-27157  oxygen partial pressure was low enough, as indicated in Reaction (5). Reaction (5) is termed as ‘active oxidation’ of SiC. Furthermore, SiO(g) would be further oxidized to form SiO2(l) in the region where the oxygen partial pressure was high enough, as indicated in Reaction (6).  SiC(s) + O2(g) → SiO(g) + CO(g)  2SiO(g) + O2(g) → 2SiO2(l)  (5)  (6)  Therefore, when the ZrB2-SiC-La2O3/SiC dual-layer coating was oxidized at 1800 °C in low-pressure airs, the oxidation products in different forms of solid (ZrO2), molten (B2O3, SiO2) and gaseous (CO, SiO) species might form. Apart from gaseous CO and SiO, the volatilization of molten B2O3 and SiO2 was also inevitable, especially in lowpressure air. These volatile species played an important role in microstructure evolution of the oxide scale as well as protection capability of the coating during oxidation at 1800 °C in low-pressure air. Therefore, to further understand the thermodynamic stability of the oxide scale, the vapor pressures of oxidation products of ZrB2 and SiC at 1800 °C under different oxygen partial pressures were calculated based on the model established by Fahrenholtz [24,25] and Yuki [26]. Finally, the volatility diagram for ZrB2-SiC at 1800 °C was plotted, as shown in Fig. 7. The relevant thermodynamic data were adopted from SGTE (Scientific Group Thermo date Europe) [27]. In the volatility diagram, only major volatile species and interest phases, such as B2O3(g), SiO(g) and SiO2(g), were involved. The vertical dash lines represented three  Fig. 6. Cross sections of the coating after oxidation at 1800 °C for 15 min in (a, b) 50, (d, e, f) 5 and (g, h, rectangle zone in figure (b).  i) 0.5 kPa air, (c) EDS element mappings of the marked  27154  \\x0c', 'Y. Ren, et al.  Ceramics International 46 (2020) 27150-27157  i  −PO2 p  i  −PO2 p  pressures of gaseous products shown in diagram far exceeded 10−3 Pa at 1800 °C under all test conditions, indicating that the volatilization of gaseous products could affect obviously the weight change as well as oxidation resistance of the coating because the volatilization of gaseous products could destroy the integrity of the surface oxide scale. It has been reported that the critical oxygen partial pressure ( a ) for the passive-to-active transition in the oxidation of SiC at 1800 °C was about 230 Pa at 1800 °C [18,29,30]. As illustrated in Fig. 7, the equilibrium oxygen pressure (PO2 ) at the SiC/SiO2 interface was 3.9 × 10−9 Pa. Therefore, at 1800 °C, the equilibrium oxygen pressure (PO2 a ) ) was always lower than the critical oxygen partial pressure ( for the passive-to-active transition in the oxidation of SiC, suggesting that in the present oxidation condition, the active oxidation of SiC in the coating occurred definitely at the coating/oxide interface. From Fig. 7, the vapor pressure of SiO(g) on the oxide surface was relevant to oxygen partial pressure, the lower the oxygen partial pressure, the higher the vapor pressure of SiO(g). In detail, at 1800 °C, the vapor pressure of SiO(g) was about 0.3, 0.1 and 0.03 Pa when the air pressure was 0.5, 5 and 50 kPa, respectively. The vapor pressure of SiO (g) in 0.5 kPa air was ten and three-folds higher than that in 50 and 5 kPa air, so SiO(g) volatilized much easily in 0.5 kPa air. Besides, it should be noted that SiO(g) formed at the coating/oxide interface due to the active oxidation of SiC could escape outwardly through the oxide scale. During such process, it would be further oxidized to form SiO2(l) following reaction (6) in the region with enough  Fig. 7. Volatility diagram for ZrB2-SiC at 1800 °C.  oxygen partial pressures applied during oxidation in this work. It can be known from Fig. 7, the vapor pressures of SiO2(g) and B2O3(g) kept constant values of about 0.08 Pa and 7.1 kPa, respectively. The former was low, but the latter was extremely high. In oxide scale, the vapor pressure of SiO(g) increased with the decrease of oxygen partial pressure, and it became highest at the interface of ZrB2-SiC and oxide scale. Generally, the weight loss induced by the volatilization of gaseous products was non-negligible when the vapor pressures of gaseous products were higher than 10−3 Pa [28]. Distinctly, the vapor  Fig. 8. Schematic illustrations of the oxidation process of the coating at 1800 °C in (a) 50, (b) 5 and (c) 0.5 kPa air.  27155  \\x0c', 'Y. Ren, et al.  Ceramics International 46 (2020) 27150-27157  high oxygen partial pressure. Apparently, the gradient of the oxygen partial pressure in the oxide scale decreased with the reduction of the oxygen partial pressure on the oxide surface. So, SiO(g) was oxidized much easily in the oxide scale in 50 kPa air.  4.2. Oxidation in 50 kPa air  −PO2 p  i  −PO2 p  For 50 kPa air, its oxygen partial pressure was 10 kPa, much higher than the critical oxygen partial pressure ( a = 230 Pa) for the passive-to-active transition in the oxidation of SiC. Therefore, the SiC outer layer was oxidized in a passive regime following reaction (1), a molten SiO2 scale formed on the coating surface. Once the oxide layer covered the whole surface of the coating, because the equilibrium oxygen pressure (PO2 = 3.9 × 10−9 Pa) at the SiC/SiO2 interface was much lower than the value of a , SiC oxidized in an active regime following reaction (5) at the SiC/SiO2 interface. The as-formed SiO(g) at the SiC/ SiO2 interface escaped outwardly through the oxide scale, a part of SiO (g) volatilized to the surrounding environment on the oxide surface, which results in the appearance of bubbles in the molten oxide scale; the other part of SiO(g) was further oxidized to form SiO2(l) in the oxide scale. Schematic illustration of the oxidation process under this condition is displayed in Fig. 8(a). It should be noted that during oxidation in 50 kPa air, the oxygen partial pressure on the oxide surface was higher compared with that in 5 and 0.5 kPa air, however, the vapor pressure of SiO(g) was lower. Besides, due to the higher air pressure on the oxide surface, the gradient of the oxygen partial pressure in the oxide scale was also higher. As a result, SiO(g) was further oxidized more easily in the oxide scale. Therefore, during oxidation in 50 kPa air, SiO(g) formed at the coating/ oxide interface was mainly oxidized again to form SiO2(l), so only minor SiO(g) volatilized and left in the form of bubbles on the surface of oxide scale. Under this condition, the coating possessed good protection capability due to the existence of SiO2(l) scale and exhibited weight gain in the whole oxidation process. When oxidation time was prolonged, the SiC outer layer was oxidized completely, the front of oxidation reaction reached the UHTCs inner layer of the coating, as shown in Fig. 8(a-2). The oxidation of the inner layer resulted in the formation of ZrO2 and B2O3. ZrO2 with high melting temperature usually acted as the skeleton of the oxide scale. Bubbles or holes were left on the surface of oxide scale due to the volatilization of B2O3. The presence of dispersed La2O3 in the oxide scale increased the viscosity of molten SiO2, so the oxygen permeability was further decreased [13,31]. Under this condition, the graphite substrate and SiC transition layer kept intact due to the protection of stable oxide scale.  4.3. Oxidation in 5 kPa air  −PO2 p  For 5 kPa air, the oxygen partial pressure is 1 kPa, also higher than the critical oxygen partial pressure ( a = 230 Pa) for the passive-toactive transition in the oxidation of SiC. Therefore, under this condition, the oxidation performance of the coating was quite similar to that in 50 kPa air, i.e. passive oxidation of the SiC outer layer occurred at the initial stage, then the SiC oxidized in active regime at the SiC/SiO2 interface and SiO(g) formed. But for this case, the vapor pressure of SiO (g) on the oxide top surface was higher and the gradient of oxygen partial pressure in the molten oxide scale was smaller compared with that in 50 kPa air. As a result, more SiO(g) volatilized and left as bubbles on the surface of oxide scale, only a small quantity of SiO(g) was oxidized to form SiO2(l). The oxidation process of the coating in 5 kPa air is shown schematically in Fig. 8(b). Moreover, due to the appearance of a great number of bubbles, and these bubbles could break, finally the integrity of SiO2 scale was deso the ZrB2-SiC-La2O3 stroyed, inner layer of the main coating was exposed into the tested environment. The ZrB2-SiC-La2O3 inner layer was oxidized to form ZrO2 and SiO2. Because the diffusion coefficient of  oxygen in ZrO2 was much higher than that in SiO2 at high/ultra-high temperatures [13], oxygen diffused through ZrO2-rich scale was more rapid than through SiO2 scale. So, as shown in Fig. 8(b-3), once bubbles blasted, the inner layer of the coating was exposed and oxidized at such local area, the oxidation rate of the coating was accelerated. Therefore, during exposure in 5 kPa air, the coating exhibited inferior protection capability due to the low integrity of the surface oxide scale and exhibited weight loss in the whole oxidation process.  4.4. Oxidation in 0.5 kPa air  −PO2 p  For 0.5 kPa air, the oxygen partial pressure is 100 Pa, which is lower than the critical value ( a = 230 Pa) for the passive-to-active transition in the oxidation of SiC at 1800 °C. In this case, the SiC outer layer of the coating was directly oxidized in an active regime following reaction (5). As illustrated in Fig. 7, the vapor pressure of SiO(g) in 0.5 kPa air was the highest one compared with the other conditions, the as-formed SiO(g) volatilized severely into ambient surrounding, which led to rapid consumption and significant weight loss of the coating. The oxidation process of the coating in 0.5 kPa air is shown schematically in Fig. 8 (c). Initially, the SiC outer layer was oxidized in an active regime, as shown in Fig. 8(c-1). As-formed SiO(g) volatilized into ambient environment immediately due to its high vapor pressure, so no condensed silicon-containing oxide formed on the coating surface. With the increase of oxidation time, the SiC outer layer was depleted and minor Mullite was left. Then, the ZrB2-SiC-La2O3 inner layer was exposed to test environment, its oxidation also produced amount of gaseous products, i.e., SiO(g), CO(g) and B2O3(g). La2O3 in the inner layer was left behind on ZrO2 scale rather than dispersing in SiO2, because Si-containing phase volatilized in the form of gas (Figs. 8(c-2)). Due to the poor sintering ability of ZrO2 and the volatilization of massive gaseous products, a great number of pores were produced in the oxide scale. Therefore, the as-formed oxide scale possessed poor oxidation resistance. Meanwhile, Mullite decomposed at ~1800 °C in low-pressure air (Reaction (7)) [32,33] and SiO2 also volatilized. The left condensed Al2O3 reacted with La2O3 to form LaAl11O18 (Reaction (8)).  Mullite(s) → 2SiO2(l) + 3Al2O3(s)  La2O3(s) + 11Al2O3(s) → 2LaAl11O18(s)  (7)  (8)  Finally, as shown in Fig. 8(c-3), porous and refractory ZrO2 based oxide scale covered on the siliconized graphite substrate. The siliconized graphite substrate would suffer from further catastrophic oxidation due to the non-protective oxide scale.  5. Conclusion  The oxidation behaviors of a dual-layer coating with the inner layer of ZrB2-SiC-La2O3 and the outer layer of SiC at 1800 °C in low-pressure air of 50, 5 and 0.5 kPa were investigated. Meanwhile, the microstructure evolution of the as-prepared coating during the oxidation was examined. The following conclusions can be drawn:  (1) Air pressure exerts a significant effect on the oxidation behavior of the as-prepared coating, mainly by affecting the volatilization behavior of SiO(g) product. The oxidation resistance of the coating decreased with the reduction of air pressure. (2) A continuous and protective oxide scale composed of ZrO2 and SiO2 with La dispersed was formed on the coating surface in 50 kPa air. With the decrease of air pressure, porous ZrO2 oxide layer was left on the coating surface due to the severe surface volatilization of SiO (g). (3) The active oxidation of SiC and the volatilization of SiO(g) are challenges for the service life of UHTC coating containing SiC in low-pressure air, which should be taken into account during the design process of a coating system.  27156  \\x0c', 'Y. Ren, et al.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgments  This work is financially supported by National Natural Foundation of China (Grant No. 51571205, 51571203).  Science  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2020.07.195 .  References  [2]  [1]  [4]  J.E. Sheehan, K.W. Buesking, B.J. Sullivan, Carbon-carbon composites, Annu. Rev. Mater. Sci. 24 (1994) 19-44, https://doi.org/10.1146/annurev.ms.24.080194. 000315. P. Crocker, B. 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  "_id": 187,
  "PDF": "Oxidation resistance and strength retention of ZrB2–SiC ceramics.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 2387-2395  Oxidation resistance and strength retention of ZrB2-SiC ceramics  Wei-Ming Guo, Guo-Jun Zhang  ∗  State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai  200050, China  Available online 4 March 2010  Abstract     Oxidation behavior and effect of oxidation on the room-temperature flexural strength were investigated for ZrB2 -10 vol% SiC (ZB10S) and ZrB2 -30 vol% SiC (ZB30S) in air at 1500 C with times ranging from 0.5 h to 10 h. The oxide scale of both ZB10S and ZB30S was composed of an outer glassy layer and an inner extended SiC-depleted layer. The changes in weight gain, glass layer thickness, and extended SiC-depleted layer thickness with oxidation were measured. Analysis suggested that the extended SiC-depleted layer was most indicative for evaluating the oxidation resistance. Compared to the ZB10S, the improved oxidation resistance in ZB30S was attributed to the viscosity increase of glassy layer and the lower number of ZrO2 inclusions in the glassy layer. Because of the healing of surface flaws by the glassy layer, the strength increased significantly by 110% for ZB10S and by 130% for ZB30S after oxidation for 0.5 h. © 2010 Elsevier Ltd. All rights reserved.  Keywords: B. Microstructure; C. Strength; ZrB2 -SiC; Oxidation; Kinetics  1.  Introduction  Among ultra high-temperature ceramics (UHTCs), ZrB2 has a desirable combination of low theoretical density, high melting temperature and thermal conductivity, which makes it attractive for use in thermal protection systems and scramjet engine components for hypersonic flight vehicles.1 When ZrB2 is exposed to air at high temperatures, it reacts with O2 to form ZrO2 and B2O3 :  ZrB2 + 5 2  O2 → ZrO2 + B2O3 (l)     Due to high vapor pressure, 1100 C2 : B2O3 (l) → B2O3 (g)  (1)  the B2O3  evaporates  above  (2)     The removal of B2O3 by vaporization leaves behind a porous ZrO2 scale, which results in the rapid linear oxidation kinetics above 1400 C.3 Thus, the high-temperature applications of monolithic ZrB2 will be limited by its poor oxidation resistance. The addition of SiC has been reported to improve the oxidation resistance of ZrB2 .3-5 Above 1100 C, SiC reacts with O2     ∗  Corresponding author. Tel.: +86 21 52411080; fax: +86 21 52413122.  E-mail address: gjzhang@mail.sic.ac.cn (G.-J. Zhang).  0955-2219/$ - see front matter © 2010 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2010.01.028  according to the following reaction:  SiC + 3 2  O2 → SiO2 (l) + CO(g)  (3)     The SiO2 and B2O3 form a borosilicate liquid, which covers the exposed surfaces. As temperature increases, B2O3 is continuously removed from the borosilicate liquid, leading to the formation of a SiO2 -rich glassy layer. Because SiO2 is significantly less volatile and more viscous than B2O3 , the SiO2 rich layer provides effective oxidation protection for ZrB2-SiC above 1100 C.5 In the past 5 years, groups in the United States, Italy and China have investigated the oxidation behavior of ZrB2-SiC ceramics.2-12 Fahrenholtz et al. have studied the structure of oxide scales on ZrB2-SiC ceramics after oxidation at temperaC.3,5,6 They indicated that the typical scale is tures up to 1500 composed of three layers: (1) a SiO2 -rich glassy layer; (2) a thin ZrO2 -SiO2 layer; (3) a SiC-depleted layer.5,6 The development of the layered structure was analyzed with the aid of a thermodynamic model that involved volatility diagrams for ZrB2 and SiC.6 The model suggested that the formation of the SiCdepleted layer was due to the active oxidation of SiC under the oxide scale.6 Similarly, Carney et al. showed that the oxide scale of ZrB2-SiC ceramics after oxidation at temperatures ranging from 1400 C to 1600 C in air was again composed of three layers, where the third inner layer was constituted by a ZrO2           \\x0c', '2388  W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395     matrix enclosing partially oxidized ZrB2 with Si-C-B-O glass inclusions.12 Oxidation studies conducted by Karlsdottir et al. have focused on the surface features of the outermost glassy layer of ZrB2-SiC ceramics that were oxidized at 1550 C in air.7-9 Island-inlagoon patterns were observed on oxide scales, consisting of a central ZrO2 “island” in a SiO2 -rich “lagoon”. These oxidescale features were called convection cells because of their role in transporting the B2O3 -rich liquid to the surface where the B2O3 was lost by evaporation, and ZrO2 precipitated from the remaining SiO2 -rich liquid. Additionally, the effect of SiC content on the formation of convection cell features was studied, showing that fewer convection cells formed and they were less uniformly distributed for ZrB2 -SiC with higher SiC content.9 The strength of non-oxide ceramics is affected by the oxidation process. To better drive high-temperature engineering applications of non-oxide ceramics, an understanding of the effects of oxidation on the room-temperature flexural strength of the material is essential. Considerable studies have reported strength retention for SiC and Si3N4 ceramics after oxidation,13,14 but only limited experimental work involves strength retention of ZrB2 -based ceramics after oxidation.15,16 Guo et al. compared the flexural strength of ZrB2-SiC composites with nano-sized or micro-sized SiC particles before and after oxidation in dry air at 1400 C for 10 h.16 After oxidation, the flexural strength increased for ZrB2-SiC containing nanosized SiC particles, whereas the strength of the ZrB2-SiC with micro-sized SiC particles decreased.16 In the present work, the oxidation behavior was studied for ZrB2-SiC ceramics containing 10 vol% and 30 vol% SiC in air at 1500 C for 0.5-10 h. Firstly, the microstructural features after oxidation were reported and discussed. Then oxidation kinetics was analyzed on the basis of weight gain, glass layer thickness and the extended SiC-depleted layer thickness. This work attempts to determine the most suitable parameter for oxidation resistance evaluation. In addition, the effect of oxidation time on the room-temperature flexural strength is also studied.        2. Experimental procedure  The raw materials used were ZrB2 (D10 = 0.46 \\u242em, D90 = 22.4 \\u242em, 98.5%, Gongyi Sanxing Ceramics Materials Co. Ltd., Gongyi, China) and SiC (D10 = 0.18 \\u242em, D90 = 1.0 \\u242em, 98.5%, Changle Xinyuan Carborundum Micropowder Co. Ltd., Changle, China). In this paper, ZrB2-SiC composites containing 10 vol% and 30 vol% SiC are referred to as ZB10S and ZB30S. The starting powder mixtures were ball milled for 8 h in acetone using Si3N4 balls in a planetary ball mill in nylon containers, and dried by rotary evaporation. Powder compacts were hot pressed at 2000 C for 60 min under a pressure of 30 MPa. At temperatures below 1650 C, the furnace was heated under vacuum. Above 1650 ing argon gas. A heating rate of 10 C, the atmosphere was switched to flowC/min was used between room temperature and 2000 of 20 C. The furnace was cooled at a rate C/min. Bars with dimensions of 2 mm × 2.5 mm × 39 mm were cut from the hot-pressed billets and ground to a 5 \\u242em surface finish                    Fig. 1. Schematic of the specimen configuration for oxidation testing. Four sites  (M and N on the top; O and P on the bottom) were at the one-third of length of  specimen.        for oxidation testing. Oxidation studies were conducted in a box furnace. The samples were placed on alumina plate with minimal contact area. For evaluating whether the formed glassy layer is fluid at the set oxidation temperature, the specimens were tilted as shown in Fig. 1. The oxidation of the hot-pressed materials was conducted at 1500 C for 0.5 h, 3 h and 10 h in stagnant air and the heating rate was 10 C/min. Weights before and after oxidation were measured using a balance with 0.1 mg precision. The weight change results were an average of three bars. The hot-pressed and oxidized specimens were cut and polished for microstructure observation. Microstructures were characterized using scanning electron microscopy (SEM) imaging in an electron probe microanalyzer (JEOL JXA-8100F, Japan) along with energy-dispersive spectroscopy (EDS, Oxford INCA energy) for chemical analysis. Flexural strength of the as sintered bars and oxidized bars was determined via threepoint bending with a span of 30 mm at a crosshead speed of 0.5 mm/min. For each composition and oxidation time, three specimens were tested.  3. Results  3.1. Microstructures of hot-pressed ZrB2-SiC ceramics  Fig. 2 shows the microstructures of full dense ZB10S and ZB30S after hot pressing. Some pits (similar appearance to pores) were apparent in the micrographs as a consequence of particle pullout during polishing. The SiC particles (black phase) were well-distributed within the ZrB2 (gray phase) matrix, and were primarily located at ZrB2 grain junctions. The SiC particles tended to be isolated in ZB10S, but had a more interconnected appearance in the ZB30S. The ZrB2 grain size of ZB10S was 3.3 \\u242em, whereas that of ZB30S decreased to 2.2 \\u242em. This showed that SiC particles inhibited the grain growth of ZrB2 during hot pressing and the effect was stronger as the volume fraction of SiC increased.  3.2. Oxidation behavior and kinetics of ZrB2-SiC ceramics  Fig. 3 presents mass gain per unit surface area as a function of oxidation time at 1500 C for ZB10S and ZB30S. With increasing oxidation time, mass gain increased for both ZB10S and ZB30S. At identical oxidation time, the mass gain of ZB10S     \\x0c', 'W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  2389  Fig. 2. Microstructures of hot-pressed ZrB2 -SiC composites: (a) ZB10S and (b) ZB30S.     was higher than that of ZB30S. For example, the mass gain of ZB10S after oxidation for 10 h was 17.2 mg/cm2 , whereas the mass gain of ZB30S was only 6.3 mg/cm2 . To understand the effect of oxidation at different points on the sample, four sites (M, N, O and P in Fig. 1) on the top and bottom of a bar were chosen for microstructural analysis. The glassy layer thickness at M was similar with that of N, and the glassy layer thickness at O was close to that at P. This implies that the inclination of about 20 has no effect on oxidation. However, the glassy layer thickness of the top of the sample was different from that of the bottom. The glassy layer thickness for the top and bottom of oxidized ZB30S was measured by SEM imaging (Fig. 4). As expected, the glass layer thickness increased with oxidation time on the top of the sample. However, on the bottom, the glass layer thickness first increased and then decreased slightly with oxidation time. Additionally, the glassy layer on the bottom was thicker than that on the top after oxidation for 0.5 h and 3 h. To better understand the oxidation behavior, the microstructure of the top of the bars was further analyzed. Figs. 5 and 6 show the backscattered electron (BSE) images and compositional maps of top of ZB10S ceramics after oxidation at 1500 C for 0.5 h and 10 h. Figs. 7 and 8 show the BSE images and com       positional maps of top of ZB30S ceramics after oxidation at 1500 C for 0.5 h and 10 h. Even with 10 vol% SiC, an outer glassy layer was formed. However, the glassy layers on ZB10S and ZB30S had different microstructural features. ZrO2 inclusions were observed in the glassy layer on ZB10S, whereas very few ZrO2 inclusions were found in the glassy layer on ZB30S. With increasing oxidation time, the size and amount of ZrO2 inclusions increased. After oxidation for 0.5 h, the size of ZrO2 inclusions on ZB10S was in the range of 1-2 \\u242em. When the oxidation time increased to 10 h, the ZrO2 inclusions grew to 5-15 \\u242em and became rod-shaped. Despite differences in the oxide layer thickness, both ZB10S and ZB30S had similar layered structures in their oxide scales. According to previous studies,5,6 the oxide scale of ZrB2 -SiC can be divided into three layers, as mentioned before. However, the ZrO2-SiO2 layer was usually very thin, so it was analyzed along with the SiC-depleted layer as one layer called the extended SiC-depleted layer. Therefore, the oxide scale for ZB10S and ZB30S consisted of two layers: (1) an outer glassy layer and (2) an inner extended SiC-depleted layer. Fig. 9 shows the thickness of the glassy layer and the extended SiC-depleted layer as a function of oxidation time for the ZB10S and ZB30S specimens at 1500 C.     Fig. 3. Mass gain per unit surface area, w, as a function of oxidation time, t, for  Fig. 4. Glassy layer thickness, H, on the top and bottom of ZB30S ceramics as     ZB10S and ZB30S at 1500     C.  a function of oxidation time, t, at 1500  C.  \\x0c', '2390  W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  Fig. 5. Backscattered electron image (a) and compositional maps for (b) Si, (c) O, and (d) Zr of a polished section of ZB10S ceramic after oxidation at 1500     C for  0.5 h.  Fig. 6. Backscattered electron image (a) and compositional maps for (b) Si, (c) O, and (d) Zr of a polished section of ZB10S after oxidation at 1500     C for 10 h.  \\x0c', 'W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  2391  Fig. 7. Backscattered electron image (a) and compositional maps for (b) Si, (c) O, and (d) Zr of a polished section of ZB30S after oxidation at 1500     C for 0.5 h.  Fig. 8. Backscattered electron image (a) and compositional maps for (b) Si, (c) O, and (d) Zr of a polished section of ZB30S after oxidation at 1500     C for 10 h.  \\x0c', '2392  W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  Fig. 10. Effect of oxidation time at 1500  strength of ZB10S and ZB30S.     C on the room-temperature flexural  to the weight gain data, the oxidation of ZB10S and ZB30S followed parabolic kinetics. Using the glassy layer thickness data, the oxidation of ZB10S exhibited linear kinetics, whereas the oxidation of ZB30S still represented parabolic kinetics. Based on the extended SiC-depleted layer thickness data, the oxidation of ZB10S followed parabolic kinetics, whereas the oxidation of ZB30S followed cubic kinetics. Determination of which data are most suitable for evaluating the oxidation behavior is discussed in Section 4.3.  3.3. Strength retention of ZrB2 -SiC ceramics after oxidation  Oxidation influenced the strength of ZrB2-SiC ceramics (Fig. 10). As oxidation progressed, the retained flexural strength of ZB10S and ZB30S had similar trends. After 0.5 h, the flexural strength increased sharply, and then gradually decreased for longer times. However, the strengths of ZB10S and ZB30S after oxidation for 10 h were still higher than that of the unoxidized specimens.  Fig. 9. Thickness of the glass layer (a) and extended SiC-depleted layer (b) on     the top surface as a function of oxidation time at 1500  C.  Combining the mass gain data with the glassy layer thicknesses and the extended SiC-depleted layer thicknesses from Figs. 3 and 9, the oxidation kinetics were analyzed for ZB10S and ZB30S using a generalized power rate equation17,18 :  xn = kt  (4)  4. Discussion  where x is the change in mass or thickness, n is the exponent, k is the rate constant and t is the oxidation time. The kinetic parameters (n, k) and correlation coefficients (R) are listed in Table 1. Fitting of the kinetic data was good, as indicated by the values of the correlation coefficients (R > 0.99). According  4.1. Microstructural features of oxidized ZrB2 -SiC ceramics  The glassy layer on the top and bottom of oxidized ZS30S differed in thickness and evolution (Fig. 4). As the temperature  Table 1  Best-fitting kinetic parameters, n, k, and correlation coefficients, R,  for of  the oxidation of ZB10S and ZB30S based on weight gain, glassy layer and extended  SiC-depleted layer growth.  Sample  Weight gain  Glassy layer  ZB10S  ZB30S  n  2  2  −2 )n h −1 ) K ((mg cm  29.7  3.9  R  0.9987  0.9990  n  1  2  −1 ) K (\\u242emn h  5.4  34.6  R  0.9976  0.9984  Extended SiC-depleted layer −1 ) K (\\u242emn h 3.4 × 103 2.3 × 103  3  n  2  R  0.9998  0.9995  \\x0c', 'W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  2393  Table 2     Degree of material damage induced by oxidation in ZB10S and ZB30S at  1500  C in air, assuming uniform oxidation within a sample.  Sample  Degree of material damage (%)  ZB10S  ZB30S  1 h  7.1  2.2  3 h  16.8  3.4  10 h  30.4  5.1                    increases, the viscosity of glassy layer should decrease rapidly, allowing the glass to flow due to gravity. The flow of glass would result in a non-uniform thickness of the glassy layer, and consequently affect the oxidation behavior. A previous study showed that gravity influenced glass flow at 1400 C and 1500 C, but C.12 that the effect was greater at 1600 In the present study, SEM observation showed that the glassy layer on the top or bottom of oxidized ZS30S was uniform, indicating that the effect of gravity was small for an angle of 20 . So, in the present work, the effect of glass flow induced by the gravity on oxidation was ignored in the present analysis. The obvious difference in the glassy layer on the top and bottom of oxidized ZS30S was mainly attributed to the different local partial pressure. As shown in Fig. 1, the space between the bottom of the sample and the alumina plate was narrow. As the temperature reached 1500 C, the vapor pressure of B2O3 increased substantially, leading to its evaporation. The narrow gap between the sample bottom and alumina plate may have allowed B2O3 vapor to accumulate, which would suppress its evaporation. As a result, the glassy layer on bottom could remain thicker than the top for shorter times (0.5 h and 3 h). As oxidation progressed, the oxygen partial pressure in the gap between the sample bottom and alumina plate decreased as oxygen was consumed and gaseous oxidation products (CO and B2O3 ) filled the gap. Due to the stagnant atmosphere, oxygen flow on the bottom surface might have been constricted and this could be the reason why the glass layer thickness did not increase on the sample bottom for longer times (10 h). ZrO2 inclusions were observed in the glassy layer of ZB10S, but not in the glassy layer of ZB30S (Figs. 5-8). During oxidation at 1500 C, B2O3 was removed from the surface of the glassy layer by evaporation. Since, oxidation is a dynamic process, newly formed B2O3 is expected to continuously dissolve into the glassy layer. So, the glassy layer should contain some B2O3 at any time during oxidation and any point through the glassy layer. Because ZB10S contained 90 vol% ZrB2 , oxidation of ZB10S would produce more B2O3 in comparison to ZB30S, which contained 70 vol% ZrB2 . So, the B2O3 concentration in glassy layer of ZB10S was higher than that of ZB30S. Based on the calculated isothermal section at 1500 C of the ZrO2-SiO2-B2O3 phase diagram,8 increasing the B2O3 content of a SiO2-B2O3 liquid increases the ZrO2 solubility in the resulting ZrO2-SiO2-B2O3 (BSZ) liquid. Combining the above analysis with the calculated phase diagram of the ternary ZrO2-SiO2-B2O3 , it was concluded that the glassy layer of ZB10S dissolved more ZrO2 compared with that of ZB30S. Karlsdottir et al. also showed that ZrB2-20 vol% SiC and especially the ZrB2 -30 vol% SiC     had less dissolved ZrO2 compared with the ZrB2 -15 vol% SiC because of the lower B2O3 content in the glassy layers.9 In the oxidized ZB10S, the ZrO2 first dissolved into the SiO2 -B2O3 liquid to form a BSZ liquid at the inner side of the glassy layer. Then the BSZ liquid would flow toward the top of glassy layer. When the B2O3 was lost by evaporation at the outer surface, ZrO2 precipitated from the BSZ liquid. As oxidation progressed, rod-shaped ZrO2 inclusions 5-15 \\u242em long grew in the glassy layer of ZB10S.  4.2. Oxidation resistance evaluation of ZrB2-SiC ceramics  From Table 1, the oxidation of ZrB2-SiC ceramics had different kinetic characteristics based on whether weight gain, glassy layer thickness, or the extended SiC-depleted layer thickness was considered. The question is which is the most suitable for evaluating the oxidation resistance. The criteria for selection should include: (1) the degree of material damage induced by oxidation and (2) a reasonable comparison of the oxidation resistance of ZrB2-SiC with different SiC contents. The best analysis has to satisfy the two criteria simultaneously. As shown in Fig. 4, the oxidation on the top and bottom surfaces of the samples was different. This evidences showed that the local environment was heterogeneous. In addition, evaporation of B2O3 in oxidized ZrB2-SiC should depend on the glassy layer composition, thickness and viscosity. Therefore, even for samples with uniform oxidation, it is not correct to compare the oxidation resistance of ZrB2-SiC with different SiC contents using weight gain data alone. Similarly, due to the uncertainty due of the evaporation of B2O3 , it is not correct to evaluate the degree of material damage and compare the oxidation resistance of ZrB2-SiC with different SiC contents based only on the thickness of the glassy layer. For example, after oxidation for 0.5 h, the glassy layer of the ZB10S was close to that of the ZB30S, whereas the extended SiC-depleted layer of the ZB10S was much extended than that of the ZB30S (Fig. 9). With the assumption that ZrB2 and ZrO2 have a theoretical density, 1 unit volume of ZrB2 upon oxidation only produced 1.1 unit volumes of ZrO2 . Because the glassy layer covered the outermost surface of the sample, the thickness of the extended SiC-depleted layer was approximately the thickness of the material damage induced by oxidation. So, it was suitable to evaluate the degree of materials damage induced by oxidation and compare the oxidation resistance of ZrB2 -SiC with different SiC contents using the extended SiC-depleted layer thickness. Now, assuming that a uniform oxidation occurs on all the surfaces of the specimen, the degree of damage of ZB10S and ZB30S could be estimated from the extended SiC-depleted layer thickness, as shown in Table 2. After oxidation for 10 h, the degree of material damage of ZB10S was about 30.4%, whereas that of ZB30S was only about 5.1%. By this measure, the degree of material damage in ZB30S after 10 h was less than in ZB10S after 0.5 h. Based on factors discussed above, the extended SiC-depleted layer can be considered the most suitable for evaluating the oxidation resistance of ZrB2-SiC ceramics.  \\x0c', '2394  W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  4.3. Oxidation kinetics and mechanism of ZrB2-SiC ceramics  5. Summary     Based on the extended SiC-depleted layer thickness, the oxidation of the ZB10S followed parabolic kinetics, whereas the oxidation of the ZB30S exhibited cubic kinetics. Generally, the oxidation of ceramics with a protective oxide layer exhibits parabolic kinetics (n = 2), indicating a diffusioncontrolled mechanism. However, when the oxygen diffusion rate through the protective oxide layer decreases, oxidation kinetics deviate from parabolic behavior, tending toward a logarithmic relationship with scale growth resulting in n »2.17,19 For example, oxidizing, annealing, and reoxidizing CVD SiC in situ at 1300 C, Ogbuji reported that devitrification of the oxide scale caused a decrease in the oxidation rate by a factor of about 30. Oxidation kinetics showed a strong departure from parabolic behavior, and the value of n was in the range from 6 to 10.20 Therefore, the increase in the value of n observed in the present study from 2 for ZB10S to 3 for ZB30S, indicated that the glassy layer in ZB30S was more protective than that of ZB10S. Hence, we conclude that ZB30S had better oxidation resistance. After oxidation for 0.5 h, the glassy layer of the ZB10S was similar to that of the ZB30S. Increasing oxidation time from 0.5 h to 3 h resulted in an increase in the thickness of the extended SiC-depleted layer by 58 \\u242em for the ZB10S. For comparison, thickness of the extended SiC-depleted layer increased by 7 \\u242em for the ZB30S. Therefore, the glassy layer thickness was not the main factor that influenced the oxidation resistance. As discussed in Section 4.1, the B2O3 concentration of glassy layer in the ZB30S is lower than that of ZB10S. The lower B2O3 concentration induced the higher viscosity of glassy layer of the ZB30S, which should retard oxygen diffusion.9 Moreover, the dramatic decrease in the number of ZrO2 inclusions enhanced the protective effect of the glassy layer on ZB30S. Therefore, the better oxidation resistance of ZB30S was attributed to two factors: (1) the higher viscosity and (2) fewer ZrO2 inclusions in the glassy layer.     Oxidation behavior and the effect of oxidation on the roomtemperature flexural strength were investigated for ZrB2 -SiC ceramics containing 10 vol% and 30 vol% SiC. After oxidation in air at 1500 C for 0.5-10 h, the oxide scale was composed of an outer glassy layer and an inner extended SiC-depleted layer. The glassy layer on top and bottom of ZS30S specimens differed in the thickness and growth rate, which was attributed to local differences in the environments of the top and bottom of the specimens. ZrO2 inclusions were found in the glassy layers of ZB10S, because in this specimen more B2O3 was produced during oxidation and the SiO2-B2O3 liquid could dissolve a higher amount of ZrO2 . On the contrary, almost no ZrO2 inclusions were observed in ZBS30. The changes in weight gain, glass layer thickness, and extended SiC-depleted layer thickness with oxidation were measured. According to each of these parameters, oxidation of ZrB2 -SiC ceramics showed different kinetic characteristics. The thickness of the extended SiC-depleted layer was approximately the thickness of the material damage induced by oxidation, which suggested that the extended SiC-depleted layer was the most suitable to evaluate the oxidation resistance. Based on thickness of the extended SiC-depleted layer, the oxidation of the ZB10S followed parabolic kinetics, while the oxidation of the ZB30S exhibited cubic kinetics, indicating that the oxidation resistance of ZB30S was better than that of ZB10S. The improved oxidation resistance in ZB30S was attributed to the viscosity increase of glassy layer and the lower number of ZrO2 inclusions in the glassy layer. Because of the healing of surface flaws by the outer glassy layer, oxidation promoted an increase in strength. After oxidation for the flexural strength increased significantly by 110% 0.5 h, for ZB10S and by 130% for ZB30S. With further oxidation (3 h and 10 h), the flexural strengths decreased, indicating that a new population of defects was created within a thicker glassy layer.  4.4.  Strength retention  Acknowledgments  The change in bending strength with oxidation time was simincreased by 110% for ZB10S and by 130% for ZB30S ilar in ZB10S and ZB30S. After oxidation for 0.5 h, the strength (Fig. 10). An increase in strength after oxidation have been previously observed from silicon based ceramics such as SiC and Si3N4 .13,14 This increase was attributed to the formation of a thin, dense oxide layer that could heal the surface flaws resulting from sample processing and machining. Surface flaw healing was effective only when the oxide layer was dense and very thin. When the oxide layer became thicker, the flaw healing effect was counterbalanced by the generation of new defects, either within the oxide scale or at the interface between the oxide scale and bulk materials.14 Therefore, after further oxidation for 10 h, the strength was reduced by 20% for ZB10S and ZB30S the strength was 65% higher after oxi(Fig. 10), but still, dation at 1500 C for 10 h compared to the room temperature values.     This work was financially supported by the Chinese Academy of Sciences under the Program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program), the National Natural Science Foundation of China (No. 50632070), and the Science and Technology Commission of Shanghai (No. 08520707800 and No. 09ZR1435500).  References  1. Fahrenholtz WG, Hilmas GE. Refractory diborides hafnium. J Am Ceram Soc 2007;90:1347-64.  of  zirconium and  2. Li J, Lenosky TJ, Först CJ, Yip S. Thermochemical and mechanical stabil ities of the oxide scale of ZrB2 + SiC and oxygen transport mechanisms. J Am Ceram Soc 2008;91:1475-80.  3. Rezaie A, Fahrenholtz WG, Hilmas GE. Oxidation of zirconium diboride C at a low partial pressure of oxygen. J Am Ceram  silicon carbide at 1500 Soc 2006;89:3240-5.     4. Monteverde F, Bellosi A. Oxidation of ZrB2 -based ceramics in dry air. J Electrochem Soc 2003;150:B552-559.  \\x0c', 'W.-M. Guo, G.-J. Zhang / Journal of the European Ceramic Society 30 (2010) 2387-2395  2395  5. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during the  13. Kim HW, Kim HE, Song H, Ha  J. Effect of oxidation on the  room    oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27:2495-501.  C. J Eur  temperature flexural strength of Ceram Soc 1999;82:1601-4.  reaction-bonded silicon carbides. J Am  6. Fahrenholtz WG. Thermodynamic analysis of ZrB2 -SiC oxidation: formation of a SiC-depleted region. J Am Ceram Soc 2007;90:143-8.  7. Karlsdottir SN, Halloran JW, Henderson CE. Convection patterns in liquid  oxide films on ZrB2 -SiC composites oxidized at a high temperature. J Am Ceram Soc 2007;90:2863-7.  8. Karlsdottir SN, Halloran JW, Grundy AN. Zirconia transport by liquid  convection during oxidation of zirconium diboride-silicon carbide. J Am Ceram Soc 2008;91:272-7.  9. Karlsdottir SN, Halloran JW. Oxidation of ZrB2 -SiC: content on solid and liquid oxide phase formation. 2009;92:481-6.  influence of SiC  J Am Ceram Soc  14. Park H, Kim HW, Kim HE. Oxidation and strength retention of monolithic  Si3N4 and nanocomposite Si3N4 -SiC with Yb2O3 as a sintering aid. J Am Ceram Soc 1998;81:2130-4.  15. Sciti D, Brach M, Bellosi A. Long-term oxidation behavior and mechani cal strength degradation of a pressurelessly sintered ZrB2 -MoSi2 ceramic. Scripta Mater 2005;53:1297-302.  16. Guo SQ, Yang JM, Tanaka H, Kagawa Y. Effect of thermal exposure on  strength of ZrB2 -based composites with nano-sized SiC particles. Compos Sci Technol 2008;68:3033-40.  17. Thorley M, Banks R. Kinetics and mechanism of oxidation of silicon nitride bonded silicon carbide ceramic. J Thermal Anal Calorim 1994;42:811-  10. Monteverde F, Scatteia L. Resistance to thermal shock and to oxidation of  22.  metal diborides-SiC ceramics for aerospace application. J Am Ceram Soc 2007;90:1130-8.  18. Wen CH, Wu TM, Wei WCJ. Oxidation kinetics of LaB6 in oxygen rich conditions. J Eur Ceram Soc 2004;24:3235-43.  11. Guo WM, Zhou XJ, Zhang GJ, Kan YM, Li YG, Wang PL. Effect of Si  19. Nickel KG. Multiple law modelling for the oxidation of advanced ceramics  and Zr additions on oxidation resistance of hot-pressed ZrB2 -SiC composites with polycarbosilane as a precursor at 1500 2009;471:153-6.  C. J Alloys Compd     12. Carney CM, Mogilvesky P, Parthasarathy TA. Oxidation behavior of zir conium diboride silicon carbide produced by the spark plasma sintering method. J Am Ceram Soc 2009;92:2046-52.  and a model-independent figure of merit.  In: Nickel KG, editor. Corro sion of advanced ceramics-measurement and modelling. Dordrecht (NL):  Kluwer Academic Pub.; 1994. p. 59-72.  20. Ogbuji LUJT. Effect of oxide devitrification on oxidation kinetics of SiC. J Am Ceram Soc 1997;80:1544-50.  \\x0c']"
},{
  "_id": 188,
  "PDF": "Oxidation Resistance of Fully Dense ZrB2 with SiC, TaB2, and TaSi2 Additives.pdf",
  "Text": "['Oxidation Resistance of Fully Dense ZrB2 with SiC, TaB2, and  TaSi2 Additives  Fei Peng and Robert F. Speyer  w  School of Materials Science and Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332-0245  Specimens of ZrB2 concentrations of B4C, SiC, TaB2, and TaSi2 were pressureless-sintered and post-hot isostatic pressed to their theoretical densities. Oxidation  containing various  resistances were studied by scanning thermogravimetry over the range 11501-15501C. SiC additions improved oxidation resistance over a broadening range of temperatures with increasing SiC con tent. Tantalum additions to ZrB2-B4C-SiC in the form of TaB2 and/or TaSi2 increased oxidation resistance over the entire evaluated spectrum of temperatures. TaSi2 proved to be a more effective additive than TaB2. Silicon-containing compositions formed a glassy surface layer, covering an interior oxide layer. This in terior layer was less porous in tantalum-containing compositions.  I.  Introduction  TRANSITION metal borides including ZrB2, TiB2, TaB2, and HfB2 are of interest for their unique properties such as ultrahigh melting temperature (430001C), high hardness and conductivities.1-3  strength,  and  high  thermal  and  electrical  They are candidates for leading edges on aircrafts and re-entry  vehicles,  as well  as  for  structural parts  in high temperature  environments. Engineering of  these ceramics for these applica tions  has  focused  on  sintering  aid  additions  to  facilitate  pressureless sintering of near-net-shape parts, and additions to  increase oxidation resistance through the formation of a passive  amorphous oxide surface coating.  Historically, ZrB2 has been densified through hot pressing in order to achieve high relative densities for improved mechanical  properties. Relative density must also be high enough so that  porosity does not mitigate oxidation resistance. More recently,  pressureless  sintering methods  to achieve high relative density  have been developed. Sciti et al. found that a ZrB2-MoSi2 powder mixture sintered well, likely via a liquid-phase sintering  mechanism, and the silicon constituent facilitated improved oxidation resistance similar to the silicon from SiC additions.4  Fahrenholtz et al. have shown that high relative density can be  achieved from pressureless  sintering of  zirconium diboride  if  boron oxide particle  coatings are  removed by vacuum heat treatment before the onset of sintering, and via the use of boron  carbide additives which react with ZrO2 ZrB2, CO(g) and volatile B2O3(g)) which would otherwise hinder sintering.5,6 In other work, they found that WC additions (in impurities  (to form  troduced via milling), and a more patient sintering period (540 min at 21501C) also resulted in high relative density (B98%).7  ZrB2 oxygen to form ZrO2 and B2O3. The B2O3 scale is nonprotective  exposed to air  at  elevated temperatures  reacts with  since boria has a high vapor pressure above B12001C (boiling 18601C). Silicon carbide  point,  i.e.  1  atm vapor pressure,  is  additions have been shown to improve the oxidation resistance  of ZrB2 via formation of an amorphous borosilicate surface coating. For ZrB2120 vol% SiC, oxidation heat treatments at 12001C and below showed weight gain no less extensive than those of specimens composed of ZrB2 alone. However, above 12001C, a borosilicate coating forms which has proven to be more imperto atmospheric oxygen penetration.8,9 Given the high  meable  volatility of boron oxide,  the silica content of  the borosilicate  glass surface coating would be expected to increase with increas ing temperature. However, the B2O3 vapor pressure may be suppressed by B2O3 entering into solution with SiO2. Compositional analysis has shown that the boron content of the oxide layer after heating to 15001C for 30 min is o1 wt%.10 Opila et al. showed that a ZrB2-20 vol% SiC composition (repeated 10 \\x02) at 13271 exposed to 10 min oxidation cycles and 16271C developed protective oxide scales: 30 mm at 13271C and 150 mm at 16271C.11 Thermal cycling at 19271C resulted in an oxide layer thickness of over 1 mm. The 16271C surface oxide  coating was identified (via energy dispersive spectroscopy) to be  silica. Underneath this coating was a region of ZrO2 dispersed in silica, which in turn was above a region of ZrB2 depleted of SiC. This region was argued to have resulted from active oxidation of SiC to form SiO(g).11,12 Opeka et al.13 have suggested that formation of SiO(g) could build up to pressures exceeding ambient, facilitating rupture of the protective glass layer, result ing in a cyclic protective/nonprotective scale-forming sequence.  In low oxygen partial pressures, formed by a CO(g)/CO2(g) mixture, ZrB2 oxidized to ZrO2(s) and a volatile boron oxide, and SiC oxidized to carbon monoxide and SiO(g). This left a nonprotective (porous) ZrO2 scale.10 In a later work, Opila et al.14 found that TaSi2 additions in the form of a ZrB2-20 vol% SiC-20 vol% TaSi2 composition showed a lower oxidation rate after cyclic oxidation at 16271C  than a ZrB2-20 vol% SiC composition. resistance was related to evidence of phase separation in the amorphous surface layer. Talmy et al.15 demonstrated the pres Improved oxidation  ence of phase separation in a ZrB2-2 vol% Ta5Si3 system based on a periodic pattern of glassy and crystalline areas on the oxi dized surface. In Opila et al.’s work, the composition containing  TaSi2 showed rapid consumption as compared to ZrB2-20 vol% SiC compositions exposed to similar oxidation heat treatments at 19271C. This was attributed to melting of Ta2O5 (17851C) and/or compounds of Ta2O5 and ZrO2. A purported advantage of tantalum compound additions is that tantalum impurity can  stabilize zirconium oxide, circumventing the tetragonal/mono clinic phase transformation (whose volume change can create  fissures in the oxide scale), and potentially reduce oxygen transport through what is normally a fast anion conductor.16 Talmy et al.17 investigated additions of Cr-, Ti-, Nb-, V-, and  Ta-borides to ZrB2-25 vol% SiC, and found that all additions (all of which formed solid solutions with ZrB2 during sintering) improved oxidation resistance over the base composition,  with TaB2 that improved oxidation resistance correlated with increasing cation field strength (defined as Z/r2, where Z is  additions being  the most  effective.  It was  found  the valence  Y. Blum—contributing editor  This project was  funded by the Air Force Office of Scientific Research, Contract  FA9550-04-1-0140.  Presented at  the AFOSR Workshop on Ultra-High-Temperature Ceramic Materials  hosted by SRI International, July 23-25, 2007.  w  Author  to whom correspondence  should be addressed.  e-mail:  robert.speyer@mse.  gatech.edu  Manuscript No. 23548. Received August 3, 2007; approved February 7, 2008.  Journal  J. Am. Ceram. Soc., 91 [5] 1489 - 1494 (2008)  DOI: 10.1111/j.1551-2916.2008.02368.x  r 2008 The American Ceramic Society  1489  \\x0c', 'of  the cation and r is the ionic radius) of  the added transition  metals.  In a borosilicate or  silicate glass with transition metal  cations,  the  tendency toward liquid immiscibility is known to  increase with increasing cation field strength of  the  transition  metal. This phase separation has been argued to result in increased viscosity,18 which has been correlated to reduced oxygen diffusion rates.13  In this  investigation,  the oxidation resistances of pore-free  pressureless-sintered and post-HIPed ZrB2-based compositions are evaluated. Scanning thermogravimetry is used as a tool to  compare  the oxidation resistance  effects of various  forms of  Ta and Si additions over a range of temperatures.  II.  Experimental Procedure  Commercially available powders were used for  raw materials.  The major crystalline phase(s), grade, and supplier are listed for  each powder in Table I. The particle sizes of commercially avail able TaB2 and TaSi2 were deemed too large for pressureless sintering. Hence, sedimentation-based selection was used to obtain  finer particles: Powders were dispersed in ethanol using an ultra sonicator (FS-14 Solid State Ultrasonicator, Fisher Laboratory  Equipment Division, Pittsburgh, PA)  for 10 min. The mixture  was allowed to settle in ethanol for 1 h. The top 7 cm of fluid was  then extracted using a pipette. Based on laser particle size analysis  (ModelLS 13 320, Beckman Coulter, Fullerton, CA), decanted particles had a d50 of 1.1 mm for TaB2 and 1.7 mm for TaSi2. The decanted suspension was dried in a beaker on a hot-plate.  The compositions of synthesized powder mixtures are given in  Table II. The powder mixtures were suspended in methanol, and  mixed in a ball mill for 24 h, using B4C as media. The milled powders were then dried in static air at 751C. The powder mix tures were then ball milled again in water with dissolved poly vinyl alcohol  (PVA, Celanese Ltd., Dallas, TX), polyethylene  glycol  (PEG, Alfa Aesar, Ward Hill, MA), and Darvan 821A  (R.T. Vanderbilt Company Inc., Norwalk, CT), using B4C as media for 8 h. PVA functioned as a binder with PEG function ing as a plasticizer and Darvan 821A served as a dispersing  agent. The highly viscous suspension after this milling step was dried in an oven at 751C, and then sieved using a 60-mesh screen.  Approximately 400 mg of powder were uniaxially pressed into  cylindrical pellets using a pressure of 117 MPa, holding for 1 min.  The pellets were loaded into latex encapsulants, which were in turn  evacuated. These were cold isostatically pressed (CIP) at 345 MPa.  This was followed by a binder-removal heat treatment of 0.251C/min to 5001C, holding for 24 h under vacuum (B4 Pa).  Fifteen pellets were fabricated for each composition.  These pellets were fired in a graphite tube furnace (Model  M11, Centorr Vacuum Industries  Inc., Nashua, NH) with  flowing argon, using graphite setters. The furnace was initially evacuated to B4 Pa (roughing pump) and backfilled with argon. The typical heating schedule was heating at 501C/min to 20001C, soaking for 1 h, and then cooling at 401C/min to room temper ature. The pellets were then hot isostatic pressed (HIP, American Isostatic Press, Columbus, OH) at 18001C for 30 min at a  pressure of 207 MPa. The densities of unfired pellets were de termined from measured dimensions and mass; the densities of  pressureless  sintered and post-HIPed pellets were determined  using Archimedes’ method.  All of  the surfaces of all post-HIPed samples were ground  away using 320 grit SiC grinding paper  (Buehler, Lake Bluff,  IL), and the  resulting pellet dimensions were measured with  calipers. The oxidation behaviors were then investigated using  thermogravimetric  analysis  (TG analysis, Model  STA 409,  Netzsch, Exton, PA, with an Innovative Thermal Systems  in terface, Atlanta, GA). The samples were supported on alumina  chips, which filled an alumina crucible, to minimize the contact  between sample surfaces and alumina. Samples were exposed  to flowing air from a compressed air tank with a flow rate of  0.1 L/min. Flow rate was maintained via a mass flow controller  (Model GFC 17, Aalborg, Orangeburg, NY). Specimens were heated from room temperature at 501C/min to 9501C, 301C/min to 10701C, 101C/min to 11001C, 51C/min to 11501C, and finally 31C/min to 15501C. Displayed data are from 11501 to 15501C in  which weight-change data were artificially shifted to zero from their values at 11501C. To evaluate repeatability, 3-4 TG oxi dation heat-treatments were performed on each composition.  Crystalline phases in the samples were identified using X-ray  diffraction (XRD, Model X’Pert PRO Alpha-1, PANalytical,  the Netherlands). Scans were recorded at room temperature over a 2y range of 101-801 at a scan rate of 0.011/s. Before oxidation  Table I.  Raw Material Characteristics  Phases  Particle size  Supplier  ZrB2  ZrB2  d50 5 2.20 mm  Grade B, H. C.  Starck, GmbH  B4C  stoichio metric B4C  d50 5 0.8 mm  Grade HS,  H. C. Starck,  GmbH  SiC  a-SiC  d50 5 0.88 mm  Grade  8S490NDP,  Superior  Graphite,  Chicago, IL  TaB2  TaB2, Ta3B4 TaSi2  o43 mm  ESPI Metals,  Ashland, OR  TaSi2  o43 mm  Cerac Inc.,  Milwaukee, WI  Table II.  Sample Compositions  Code  ZrB2  B4C  SiC  TaB2  TaSi2  Vol%  Mol%  Vol%  Mol%  Vol%  Mol%  Vol%  Mol%  Vol%  Mol%  ZB3  90.0  91.4  10.0  8.6  0.0  0.0  0.0  0.0  0.0  0.0  ZB5  86.5  88.3  13.5  11.7  0.0  0.0  0.0  0.0  0.0  0.0  ZBS2  80.4  77.4  8.9  7.2  10.7  15.3  0.0  0.0  0.0  0.0  ZBS4  77.6  73.6  8.6  6.9  13.8  19.5  0.0  0.0  0.0  0.0  ZBS6  75.0  70.2  8.3  6.6  16.7  23.2  0.0  0.0  0.0  0.0  ZBS8  72.6  67.1  8.1  6.3  19.3  26.5  0.0  0.0  0.0  0.0  ZBS10  70.3  64.2  7.8  6.0  21.9  29.7  0.0  0.0  0.0  0.0  ZBS14  64.7  57.5  7.2  5.4  28.1  37.1  0.0  0.0  0.0  0.0  ZBS18  58.8  50.7  6.5  4.7  34.7  44.5  0.0  0.0  0.0  0.0  ZTBS1-1  65.2  59.8  7.3  5.7  20.5  27.9  7.0  6.6  0.0  0.0  ZTBS1-5  64.2  59.8  7.1  5.6  20.2  28.0  3.5  3.4  5.0  3.3  ZTBS1-9  63.3  59.8  7.0  5.6  19.9  28.0  0.0  0.0  9.8  6.6  ZTBS2-1  58.1  53.2  7.3  5.6  20.5  27.9  14.1  13.3  0.0  0.0  1490  Journal of the American Ceramic Society—Peng and Speyer  Vol. 91, No. 5  \\x0c', 'heat-treatment,  specimen  surfaces  were  ground  to  expose  specimen interior regions for XRD analysis. Microstructures of  oxidized samples were investigated using scanning electron mi croscopy (SEM, LEO 1530, Carl Zeiss SMT Inc., Thornwood,  NY). Specimen cross-sections were formed by cutting through the  cylinder axis of  the specimens with a diamond wafering blade  (Isomet 4000, Buehler, Lake Bluff, IL). Specimens were coated  with gold (sputtering for 2 min)  to form a conductive surface.  Specimens evaluated by XRD and SEM were surface-oxidized in  the TG furnace using the same schedule as previously indicated,  with the exception that they were exposed to ever-decreasing heating rates between 14301 and 14601C (B11C/min at 14601C).  III.  Results  Table III shows the green relative densities of  the pellets after  uniaxial and hydrostatic pressing (CIP), along with relative den sities after pressureless  sintering and post-HIPing. Theoretical  densities were based on the rule of mixtures. All post-HIPed  specimens were at or very near their theoretical densities.  Table IV shows X-ray diffraction results for various compo sitions after post-HIPing, and for the surfaces of  those compo sitions after exposure to TG heat-treatments in flowing air to 14601C. The presence of B4C was not detected before oxidation. For tantalum-containing specimens, ZrB2 peaks (e.g. (101), (110), and (001)) were shifted to slightly higher values of 2y, consistent  with tantalum entering into solid solution with zirconium in ZrB2.19 After oxidation, TaC was detected from XRD of specimen surfaces in tantalum-containing compositions (e.g. ZTBS1-1  in Fig. 1). Zirconia was detected in its monoclinic form.  Repeatability  among  the multiple TG traces of  the  same  composition was  found to be good; displayed TG traces are  those that were typical  for their composition. Weight changes and 11501C were B10% of between room temperature weight changes measured over the range 11501-15501C.  the  Figure 2 shows the mass gain of ZrB2 containing two concentrations of the sintering aid B4C. Mass change was linear to B13501C for the two concentrations, above which mass gain  accelerated. The B4C content did not appear to have a significant influence on oxidation resistance. The figure also shows the  mass loss of B2O3 powder, heated under the same schedule. Figure 3 shows the mass changes for ZrB2-B4C with varying concentrations of SiC additions. For the lowest concentration  of  SiC (ZBS2), mass  increase was  attenuated  compared  to  the  compositions  in Fig. 2 with no SiC additions. With SiC  additions increasing up to 13.8 vol% SiC (ZBS4), weight gain  continuously increased, accelerating over the temperature range of B14001-15001C. For SiC contents of 19.3 vol% and above  (ZBS8-ZBS18),  temperature spans of mass  loss are apparent.  These are indicated in the figure by horizontal bars.  Figure 4 shows the effect of tantalum additions, both in the  form of TaB2 and TaSi2. Weight gain was attenuated when TaSi2, or a combination of TaSi2 and TaB2 were used, as compared with when TaB2 was used alone. Figure 5 shows a reduced extent of mass gain when both ZrB2 and TaB2 were used together, as compared to when ZrB2 was used alone, for samples of similar SiC content.  Figure 6 compares the lowest mass-gain compositions (after heating to 15501C) of the different categories. SiC addition decreased the rate of mass gain starting at B12901C and there loss of 13251-14101C. SiC addition along  was a range mass  with combined TaB2 and TaSi2 position with the lowest mass gain over the entire spectrum of  additions  formed the  com evaluated temperatures.  Figure 7 shows a cross-sectional view near the surface of ZB3, heated to 14601C. Based on Energy-dispersive lighter-shaded B40-mm-thick  spectrometric  analyses,  the  band  contained  oxygen, while interior  regions did not. Boron is  too light an  element to be detected by the available EDS.  Table III.  Green, Sintered, and Post-HIPed Relative  Densities  Code  Green relative  density (%)  Sintered relative  density (%)  Post-HIPed relative  density (%)  ZB3  60.5170.66 61.9170.48 62.6270.65 63.0970.57 62.8570.63 63.5970.71 64.2770.57 65.6770.23 64.9170.88 64.6170.81 63.9270.67 63.1870.68 64.2471.25  94.50  98.9270.26 99.0670.35  ZB5  93.51  ZBS2  95.75  100  ZBS4  95.43  100  ZBS6  96.15  100  ZBS8  96.77  100  ZBS10  96.67  100  ZBS14  94.77  100  ZBS18  95.67  100  ZTBS1-1  96.01  100  ZTBS1-5  95.27  100  ZTBS1-9  94.47  100  ZTBS2-1  97.98  100  Sintered relative densities were measured for only a sampling of each compo sitions; hence, no standard deviations are indicated.  Table IV.  XRD-Identified Phases  Code  Post-HIPed phases  14601C oxidized  surface phases  ZB3  ZrB2 ZrB2 ZrB2 ZrB2, SiC ZrB2, SiC ZrB2, SiC ZrB2, SiC ZrB2, SiC ZrB2, SiC ZrB2-TaB2 (ss), SiC  ZrO2 ZrO2 ZrO2 ZrO2, ZrB2(tr) ZrO2, ZrB2(tr) ZrO2, ZrB2(tr) ZrO2, ZrB2(tr) ZrO2, ZrB2 ZrO2, ZrB2 ZrB2-TaB2 (ss), ZrO2, TaC, TaO(tr) ZrB2-TaB2 (ss), ZrO2, TaC, TaO(tr) ZrB2-TaB2 (ss), ZrO2, TaC, TaO(tr) ZrB2-TaB2 (ss), ZrO2, TaC  ZB5  ZBS2  ZBS4  ZBS6  ZBS8  ZBS10  ZBS14  ZBS18  ZTBS1-1  ZTBS1-5  ZrB2-TaB2 (ss), SiC, TaSi2(tr) ZrB2-TaB2 (ss), SiC, TaSi2(tr) ZrB2-TaB2 (ss), SiC  ZTBS1-9  ZTBS2-1  ss, solid solution ; tr,  indicates trace quantities.  20  30  40 50 60 Scattering angle (Degrees 2θ)  70  80  I  n  t  e  n  s  t i  y  (  C  n u o  t  s  )  Fig. 1. X-ray diffraction pattern of ZTBS1-1 after heat treatment 14601C. z, ZrO2; c, TaC; o, TaO; b, ZrB2-TaB2 (ss), u, unknown.  to  May 2008  Oxidation Resistance of Fully Dense ZrB2  1491    \\x0c', 'Figure 8 shows three layers in the cross-section of oxidized ZBS18. A 5-mm surface  layer with a glassy appearance  con tained silicon and oxygen, but was devoid of zirconium. A second porous layer of B20-mm thickness contained Zr, Si, O, and  trace carbon. Beneath these layers was a well-densified matrix  containing Zr and Si, but no oxygen. Figure 9 similarly shows  three distinct regions in the cross-section of oxidized ZBTS1-1. A B5-mm glassy surface coating contained silicon, oxygen, and  only a trace quantity of zirconium. There was no positive iden tification of tantalum in this layer (see Section IV). A second layer of B20-mm-thick contained zirconium,  tantalum, silicon,  carbon, and oxygen. This region appeared less porous than the  corresponding second layer  in ZBS18. The  specimen interior  contained Zr, Ta, and Si, with no oxygen detected.  IV.  Discussion  Boron carbide functioned as an effective sintering aid, facilitat ing sintered relative densities in excess of 93.5% for ZrB2 alone, and with various SiC, TaB2, and TaSi2 additions. This exceeded the closed-porosity requirement for post-HIPing to be effec tive.20 This in turn eliminated porosity as a factor in evaluation  of oxidation resistance.  Although  evaluating  oxidation  resistance  under  constant  heating rates does not  facilitate modeling of oxidation kinetics  at specific temperatures,  it  is a useful comparison tool  for the  oxidation resistances of various compositions over a spectrum of  temperatures. The use of B4C sintering aid appears harm to the oxidation resistance of ZrB2 (Fig. 2). Even at lowest evaluated temperatures, B2O3 shows volatility. However, even if the oxidation of ZrB2 resulted in strictly B2O3 vapor rather than condensed phase, a weight gain would still be ex to do no  the  pected because of  the zirconia which forms (Table V). Signifi cantly more weight is gained if the boria forms as a condensed phase. The relatively linear mass increase up to B13501C rep resents a convolution of an increasing thickness of B2O3 and ZrO2-B2O3 layers, through which oxygen must diffuse (slowing oxidation), with an increased vapor pressure of B2O3, which would remove the B2O3 semi-protective layer, accelerating oxidation. After heating to 14601C, there is no visually apparent  amorphous layer, only a zirconia scale (Fig. 7).  Slopes of mass change with temperature were the same for all (except ZBS18) up to B13501C.  SiC-containing compositions  M  a  s s  c  h  a  g n  e  (  m  g  /  c  m  )  M  a  s s  c  h  a  g n  e  (  m  g  )  0 1150  2  4  6  8  10  12  14  16  −16 1550  −14  −12  −10  − 8  − 6  − 4  − 2  0  1200  1250  1300 1350 1400 Temperature (° C)  1450  1500  ZB3  ZB5  B O  Fig. 2.  Scanning  thermogravimetric  (TG)  analysis  (31C/min)  of  ZrB2-B4C specimens. ZB3 contains 10.0 vol% B4C and ZB5 contains 13.5 vol% B4C. Superimposed is a portion of a TG trace of B2O3 powder, showing an increased rate of volatilization with increasing  temperature.  ZBS2 ZBS4 ZBS8 ZBS10 ZBS14 ZBS18  ZBS8 ZBS10 ZBS14 ZBS18  M  a  s s  c  h  a  g n  e  (  m  g  /  c  m  )  1300 1350 1400 Temperature (°C)  0 1150  2  4  6  8  10  1200  1250  1450  1500  1550  Fig. 3.  Thermogravimetric  analysis  of  ZrB2-B4C amounts of SiC. Volume percentages of SiC increased from  specimens  with  varying  10.7 for ZBS2 to 34.7 for ZBS18. Horizontal bars indicate temperature  spans of mass loss.  0 1150  0.5  1  1.5  2  2.5  3  3.5  4  1200  1250  1300 1350 1400 Temperature (° C)  1450  1500  1550  ZTBS1-1 ZTBS1-5 ZTBS1-9  M  a  s s  c  h  a  g n  e  (  m  g  /  c  m  )  Fig. 4.  Thermogravimetric analysis of ZrB2-B4C-SiC specimens with TaB2 and TaSi2 additions. ZTBS1-1 has only TaB2, ZTBS1-9 has only TaSi2, and ZTBS1-5 has both additives.  0 1150  1  2  3  4  5  6  7  1200  1250  1300 1350 1400 Temperature (° C)  1450  1500  1550  ZBS10 ZTBS2-1  M  a  s s  c  h  a  g n  e  (  m  g  /  c  m  )  Fig. 5.  Effect of substitution of TaB2 for ZrB2: ZBS10 contains only ZrB2. ZTBS2-1 contains a mixture of ZrB2 and TaB2. Concentrations of B4C and SiC were held approximately constant.  1492  Journal of the American Ceramic Society—Peng and Speyer  Vol. 91, No. 5                      \\x0c', 'This is consistent with the work of others,8,9 which shows that  over this lower temperature range, SiC does not oxidize to con tribute to forming a protective layer. For higher concentrations of  SiC (i.e. ZBS8, ZBS10, ZBS14, ZBS18), spans of mass loss are  seen;  in these temperature ranges, the net rate of mass gain from  oxidation was less than the rate of mass loss from volatilization of  B2O3 from the amorphous borosilicate surface layer. At temperatures above these ranges, mass gain resumed, indicating the  dominance of accelerating oxygen diffusion through the amor phous silica surface layer. Behind this surface layer, a porous layer  formed (Fig. 8). This could have formed via capillary extraction of  the silica from oxidized SiC to the amorphous silica surface layer, or from formation of SiO(g), as indicated in other work.11,12 For compositions with tantalum additions, during the oxida tion of SiC beneath the silica layer, the liberated carbon formed  TaC, presumably by  SiCðsÞ þ TaB2ðsÞ þ 2O2ðgÞ ! TaCðsÞ þ SiOðgÞ þ B2O3ðl Þ ; 1700 K ¼\\x007000 kJ=mol  DG\\x0e  This reaction must have consumed oxygen in these regions to  the point at which further oxidation of TaC to form Ta2O5 did not occur in the time frame of the oxidation heat-treatment.  Additions of TaSi2 than TaB2 additions (Fig. 4), though both additives were helpful. This would be expected since oxidation of both Ta and Si  resulted in greater oxidation resistance  contributes to protective species—SiO2 which forms the protective surface glass, and Ta2O5 which has been argued to facilitate phase separation (with associated viscosity increase and oxygen in the amorphous layer.14,17 This result  diffusivity decrease)  is  even more striking in light of  the fact  that a significant weight  gain should occur from the oxidation of TaSi2 (Table V). For approximately the same SiC content, tantalum additions (in the  form of TaB2) resulted in improved oxidation resistance over the entire evaluated temperature span (11501-15501C, Fig. 5).  The oxidized layer beneath the amorphous surface layer (Fig. 9)  appears less porous than that observed in the ZBS series (Fig. 8).  This may be the result of the Ta2O5-SiO2 liquid phase having a higher viscosity, leaving it less vulnerable to capillary extraction  to the amorphous surface layer. This second oxide layer would  then likely be more protective of  the underlying diboride than  the analogous layer in the ZBS series.  EDS results did not definitively show the presence or ab sence  of  tantalum in  the  amorphous  surface  coatings  of  ZBTS1-1, ZBTS1-5, and ZBTS1-9; however, there was positive  0 1150  2  4  6  8  10  12  14  16  1200  1250  1300 1350 1400 Temperature (° C)  1450  1500  1550  ZB5 ZBS18 ZTBS1-5  M  a  s s  c  h  a  g n  e  (  m  g  /  c  m  )  Fig. 6.  Comparison  of  various  compositions: ZB5  has  no  silicon bearing constituents. ZBS18  contains 34.7  vol% SiC, and ZTBS1-5  contains SiC, TaB2 and TaSi2.  10 µm  Zr, O, C  Zr, O, C  Zr,C  Fig. 7.  Secondary electron scanning electron micrograph of  the cross  section of ZB3 (90 vol% ZrB2, 10 vol% B4C) heat-treated in flowing air to 14601C. Rectangles enclose regions in which energy-dispersive  spectrometric  scans were  taken. The  elements  that were detected are  indicated next to each rectangle. Similar principles apply to Figs. 8 and 9.  10 µm  Zr, Si, C  Zr, Si, O, Ctr  Si, O, Ctr  Fig. 8.  Secondary  electron scanning  electron micrograph of ZBS18  (58.8 vol% ZrB2, 6.5 vol% B4C, and 34.7 vol% SiC) thermogravimetric oxidation heat-treatments to 14601C.  exposed  to  10 µm  Zr, Ta, Si, C, O  Zr, Ta, Si, C  Zrtr, Si,Ta? C, O  Fig. 9.  Secondary electron scanning electron micrograph of ZBTS1-1  (65.2 vol% ZrB2, 7.3 vol% B4C, 20.5 vol% SiC, and 7.0 vol% TaB2) exposed to thermogravimetric oxidation heat-treatments to 14601C.  May 2008  Oxidation Resistance of Fully Dense ZrB2  1493      \\x0c', 'identification of this element in the oxidized layers below, as well  as deeper in the matrices: Positive identification of tantalum can  be made when its ‘‘L’’  lines are apparent. The ‘‘M’’  line is the  most intense for tantalum, but it overlaps with the ‘‘K’’  line of  silicon. A characteristic of the ‘‘M’’ line is a second small peak at lower KeV,21 which was observed in EDS analysis of TaB2 alone. In the case of the surface glass, only one peak corre sponding to the ‘‘M’’  line of Ta and the ‘‘K’’  line of Si was  observed.  V.  Conclusions  Pressureless  sintering followed by post-HIPing of ZrB2, taining B4C as a sintering aid, and various compositions containing SiC, TaB2, and TaSi2, resulted in compacts reaching theoretical density. This eliminated  con the  effect  of  variable  porosity on oxidation behavior.  Increasing SiC additions  im proved  oxidation  resistance  over  an  expanding  range  of  temperatures. TaB2 substitution for ZrB2 resistance as soon as a temperature was  improved oxidation  reached in which an  amorphous borosilicate protective layer formed. TaB2 and TaSi2 additions to ZrB2-B4C-SiC improved oxidation resistance over  the entire range of evaluated temperatures. Specimens containing silicon heat-treated to 14601C formed an amorphous silica  coating, with a porous oxide layer beneath. This layer appeared  less porous in Ta-containing specimens.  Acknowledgments  The  authors would like  to acknowledge Ms. Yolande Berta  for  electron  microscopy advice,  Jack Bourbonnais  for help with specimen preparation, as  well as Schenck Wiley and Lionel Vargas  for help with binder  removal heat treatments. The authors would like to express their appreciation for the helpful  suggestions and support of their contract monitor, Dr. Joan Fuller.  References  1Y. Murata and E. B. Whitney, ‘‘Densification and Wear Resistance of Ceramic  Systems: III. Tantalum Mononitride-Zirconium Diboride,’’ Am. Ceram. Soc. Bull.,  48 [7] 698-702 (1969). 2Y. Murata, ‘‘Densification and Wear Resistance of HfN-ZrB2 Compositions,’’ Am. Ceram. Soc. Bull., 52 [3] 255-9 (1973). 3X. Zhang, P. Hu, S. Meng,  J. Han,  and B. Wang,  ‘‘Microstructure  and  Mechanical Properties of ZrB2-Based Ceramics,’’ Key Eng. Mater., 312, 287-92 (2006). 4D. Sciti, M. Brach, and A. Bellosi,  ‘‘Oxidation Behavior of a Pressureless Sin tered ZrB2-MoSi2 Ceramic Composite,’’ J. Mater. Res., 20 [4] 922-30 (2005). 5S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Pressureless Densification  of Zirconium Diboride with Boron Carbide Additions,’’ J. Am. Ceram. Soc., 89 [5]  1544-50 (2006). 6W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. Zaykoski,  ‘‘Refrac tory Diborides of Zirconium and Hafnium,’’ J. Am. Ceram. Soc., 90 [5] 1347-64  (2007). 7A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Pressureless  Sintering of Zirconium Diboride,’’ J. Am. Ceram. Soc., 89 [2] 450-6 (2006). 8W. C. Tripp, H. H. Davis, and H. C. Graham,  ‘‘Effect of an SiC Addition on  the Oxidation of ZrB,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 9M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykosi, and S. J. Causey,  ‘‘Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and  Zirconium Compounds,’’ J. Eur. Ceram. Soc., 19 [12-14] 2405-14 (1999). 10A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, ‘‘Oxidation of Zirconium Diboride-Silicon Carbide at 15001C in a Low Partial Pressure of Oxygen,’’ J. Am.  Ceram. Soc., 89 [10] 3240-5 (2006). 11E. J. Opila and M. C. Halbig,  ‘‘Oxidation of ZrB2-SiC,’’ Ceram. Eng. Sci.  Proc., 22 [3] 221-8 (2001). 12A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Evolution of Structure  During the Oxidation of Zirconium Diboride-Silicon Carbide 15001C,’’ J. Eur. Ceram. Soc., 27, 2495-501 (2007). 13M. Opeka, I. Talmy, and J. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C 1 Hypersonic Aerosurfaces: Theoretical Considerations and Historical  in Air  up  to  Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 14E. J. Opila, S. Levine, and J. Lorincz,  ‘‘Oxidation of ZrB2and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39 [19]  5969-77 (2004). 15I. G. Talmy, J. A. Zaykoski, M. M. Opeka, and A. H. Smith,  ‘‘Properties of  Ceramics in the System ZrB2-Ta5Si,’’ J. Mater. Res., 21 [10] 2593-9 (2006). 16S. R. Levine and E. J. Opila ‘‘Tantalum Addition to Zirconium Diboride for  Improved Oxidation Resistance,’’ NASA/TM-2003-212483. 17I. G. Talmy, J. A. Zaykoski, M. M. Opeka, and S. Dallek, ‘‘Oxidation of ZrB2 Ceramics Modified with SiC and Group IV-VI Transition Metal Diborides,’’ Elec.  Chem. Soc. Proc., 12, 144-58 (2001). 18W. Vogel, Glass Chemistry, 2nd Edition, Springer-Verlag, New York, 1994. 19Y. Xie, T. H. Jr. Sanders, and R. F. Speyer,  ‘‘Solution-Based Synthesis and  Processing of Sub-Micron ZrB2 and ZrB2-TaB2,’’ J. Am. Ceram. Soc., Accepted for publication. 20N. Cho, Z. Bao,  (2007).  and R. F. Speyer,  ‘‘Density and Hardness-Optimized  Pressureless Sintered and Post-HIPed B4C,’’ J. Mat. Res., 20 [8] 2110-6 (2005). 21J. I. Goldstein, D. E. Newbury, P. Echlin, D. C. Joy, A. D. Jr. Romig, C. E.  Lyman, C. Fiori,  and E. Lifshin, Scanning Electron Microscopy  and X-ray  Microanalysis, 2nd Edition, Plenum Press, New York, 1992.  &  Table V.  Mass Changes from Oxidation Reactions  Reaction  Mass change  (grams) per  mole Boride/  Silicide/Carbide  ZrB2ðsÞ þ 5 2  O2ðgÞ ! ZrO2ðsÞ þ B2O3ðgÞ  10.4  TaB2ðsÞ þ 11 4  O2ðgÞ ! 1 2  Ta2O5ðsÞ þ B2O3ðgÞ  18.4  SiCðsÞ þ 2O2ðgÞ ! SiO2ðl Þ þ CO2ðgÞ  20.0  ZrB2ðsÞ þ 5 2  O2ðgÞ ! ZrO2ðsÞ þ B2O3ðl Þ  80.0  TaB2ðsÞ þ 11 4  O2ðgÞ ! 1 2  Ta2O5ðsÞ þ B2O3ðl Þ  88.0  TaSi2ðsÞþ 13 4  O2ðgÞ ! 1 2  Ta2O5ðsÞ þ 2SiO2ðlÞ  104.0  1494  Journal of the American Ceramic Society—Peng and Speyer  Vol. 91, No. 5  \\x0c']"
},{
  "_id": 189,
  "PDF": "Oxidation Resistance of Hafnium Diboride Ceramics with Additions of Silicon Carbide and Tungsten Boride or Tungsten Carbide.pdf",
  "Text": "['Oxidation Resistance of Hafnium Diboride Ceramics with Additions of  Silicon Carbide and Tungsten Boride or Tungsten Carbide  Carmen M. Carney,*,w,z,y  Triplicane A. Parthasarathy,*,z,y  and Michael K. Cinibulk*,z  zAir Force Research Laboratory, Wright-Patterson Air Force Base, Ohio 45433-7817  y  UES Inc., Dayton, Ohio 45432  Dense samples of HfB2-SiC, HfB2-SiC-WC, and HfB2-SiC- WB were prepared by field-assisted sintering. The WB and WC  additives were incorporated by solid solution into the HfB2 and the HfC that formed during sintering. Oxidation of the samples furnace oxidation between 16001 was studied using isothermal and 20001C. Sample microstructure and chemistry before and after oxidation were analyzed by scanning electron microscopy  and X-ray diffraction. The addition of WC and WB did not alter  oxidation kinetics of the baseline HfB2-SiC composition below 18001C; however, at 20001C, HfB2-SiC-WC and HfB2-SiC- WB had oxide scales that were 30% thinner than the oxide  scale of HfB2-SiC. It liquid-phase densification of  is believed that WC and WB promoted the HfO2 the path of oxygen ingress, during oxidation.  scale,  thereby reducing  I.  Introduction  TRANSITION metal borides and carbides with melting temperatures exceeding 27001C are commonly referred to as ultra-high-temperature ceramics (UHTCs). Of these materials,  hafnium diboride (HfB2) has shown the highest oxidation resistance.1-5 Although its density is higher than that of zirconium diboride (ZrB2)1-4 (10.5 vs 6.09 g/cm3) its high melting point (431001C) and high thermal conductivity (B100 W/MK  at room temperature) make it a leading candidate for extreme  environment thermal protection systems such as those found at the sharp leading edges of hypersonic vehicles.6,7  The oxidation of HfB2 proceeds according to reaction (1) where the B2O3 liquid formed is volatile and expected to be removed through evaporation above 11001C. Although HfO2 is a stable refractory with a melting point of 27501C the oxide  formed on the surface of the diboride is porous and nonprotective from further oxygen penetration.3 As  such, development  of these materials in the 1960s led to the addition of 5-30 vol%  of silicon carbide (SiC) to form a borosilicate glass upon oxidation.1,3-5 Above 11001C, SiC will oxidize by reaction (2).  Available B2O3 SiO2 borosilicate glass that covers or fills the pores of the HfO2 scale. The SiO2 glass has shown to improve oxidation resistance (o17001C) when at moderate temperatures the oxidation of pure HfB2.3,6,7 HfB2 ðcrÞ þ 5=2 O2 ðgÞ ! HfO2 ðcrÞ þ B2O3 ðl Þ  can  dissolve  in  to  form a  protective  compared with  (1)  SiCðcrÞ þ 3=2 O2 ðgÞ ! SiO2 þ COðgÞ  (2)  However, as these materials have been tested in environments  that more closely simulate atmospheric reentry and hypersonic cruise applications that include higher temperatures (417231C) and flowing air8-13 it has become apparent that there is a point  at which the protective SiO2-based scale begins to fail. Above the melting point of SiO2 (17231C) it is observed that the less viscous SiO2 will flow into the pores of the HfO2 and even flow from the sample resulting in a less protective porous outer scale.12 Recent research has focused on extending the oxidation  resistance of HfB2-SiC to these temperatures by adding additional phases.7,13-17 There are three strategies for selecting the  appropriate additive that will accomplish one or more goals of  (1) controlling the phase transformations of the HfO2 scale, (2) increasing the viscosity of the SiO2 scale, or (3) promoting densification of the HfO2 scale. Mixed results have been obtained with the addition of rare earth elements (Ta, La and Nd) and some transition metals (Ti,  Cr, and Cr) with reduced oxidation being observed for oxidation temperatures below 17001C7,15-21 and increased oxidation above 17001C.12,15 Most recently, Zhang et al.17,21 have prepared ZrB2based UHTCs with additions of WC to promote liquid-phase  sintering in the oxide scale. It was shown that  the sample con taining WC had a denser ZrO2 scale and the weight gain during oxidation was reduced at temperatures of 16001C and below.  The present paper focuses on the addition of tungsten (W) to  an HfB2-based UHTC to promote improved oxidation resistance at temperatures 416001C. The WO3-HfO2 phase diagram suggests a solid solution between the phases of at least 5 mol% temperatures up to and beyond 20001C.22 Similar  WO3 at ZrO2-WO3, HfO2 will form a liquid phase with WO3 at temperatures above 12801C.22 In the present study, W was added in  to  two phases:  (1) WC and (2) WB. Samples with W-based addi tives are  compared with a baseline  composition of HfB2-15 vol% SiC. SiC content was held constant in the samples to allow  a direct  comparison. Fifteen volume percent SiC was  chosen  based on internal experiments that showed it to provide oxida tion scales of similar total  lengths and comparable weight gains  as 20 vol% SiC at  the tested temperatures, but with a thinner 16001  SiO2 scale. The samples were 20001C. Because of  oxidized  between  and  the differences  in weight between the for mation of HfO2 and WO3 volatility of WO3, thickness was used as a measure of oxidation resistance.  and the  scale  II.  Experimental Procedure  (1)  Powder Processing  Commercially available HfB2, b-SiC, WC, and WB powders were used to prepare the samples. Table I is a list of the sup pliers, purity, and starting particle sizes. Powders were used as  received except  for the HfB2 which was premilled using Si3N4 grinding media in isopropanol for 60 h to reduce the grain size.  The weight change of the grinding media was 0.07 g, which was  0.01 wt% of  the total batch weight. Representative batches of  milled HfB2 were measured using a laser diffraction particle size  N. Jacobson—contributing editor  This work was financially supported under Contract No. FA8650-10-D-5226.  *Member, The American Ceramic Society.  w  Author  to whom correspondence  should  be  addressed.  e-mail:  carmen.carney@  wpafb.af.mil  Manuscript No. 28917. Received November 19, 2010; approved January 21, 2011.  Journal  J. Am. Ceram. Soc., 94 [8] 2600 - 2607 (2011)  DOI: 10.1111/j.1551-2916.2011.04462.x  r 2011 The American Ceramic Society  2600  \\x0c', 'analyzer (LS230, Beckman Coulter, Brea, CA) and had an average particle size of 1.3 mm (D90 5 2.2 mm). The oxygen content of the premilled HfB2 was 0.86 wt% compared with 0.58 wt% in the as received powder as measured by Leco Corporation (St.  Joseph, MI). Compositions were chosen based on the HfO2- WO3 phase diagram. The prepared samples were HfB2-15 vol% SiC (HS), HfB2-15 vol% SiC-3 vol% WC (HSWC) and HfB2- 15 vol% SiC-3 vol% WB (HSWB). The powder mixtures were  ball milled in isopropanol  for 24 h with SiC grinding media,  dried at  room temperature, and subsequently dry milled for  12 h. Typical weight loss of the SiC grinding media after milling  the additives with HfB2 was 0.2 mg (0.1 wt% of the total batch). The powders were sieved through an 80-mesh (177 mm) screen.  (2)  Sintering and Sample Preparation  The milled powders were loaded into a 25-mm-diameter graphite  die. A layer of BN and graphite foil separated the powder from  the die with the powder  in contact with the graphite foil. The  powder-filled dies were cold pressed at approximately 20 MPa  and loaded into the field-assisted sintering (FCT Systeme GmbH,  Model HPD 25-1, Rauenstein, Germany) unit. The graphite die  was wrapped in graphite felt to limit heat radiation from the die. Heating and cooling rates were 501C/min. and 32 MPa was applied while heating up to 16001C. The 32 MPa was held for the  remainder of the sintering schedule and rereleased during the free cool (below 10001C). The samples were held at 21001C for times  between 5 and 9 min. The temperature was measured by a py rometer focused on the bottom of a bore hole in the upper punch B5 mm from the surface of the powder. Densification was mon itored by tracking the movement of the pistons.  Sintered samples were cut with a wire electron discharge 2.5 mm \\x02 2.0 mm \\x02 9.0 mm rectangles machine into and polished using diamond slurry to a 1 mm finish on all six sides  using an autopolisher to ensure samples with parallel sides and  uniform sizes.  (3)  Oxidation Exposure and Sample Characterization  Polished samples were heated in a zirconia element furnace (ZrF-25: Shinagawa Refractories Co., Tokyo, Japan) to 16001, 18001, or 20001C and held at temperature for 30 min. The heating and cooling rates were 51C/min., which were limited by the  furnace. Samples were placed on a concave ZrO2 crucible to limit contact between the sample and ZrO2. Oxidized samples were polished in cross section perpendicular to the bottom (2.5 mm \\x02 9.0 mm side facing the crucible) of the sample to a 1 mm finish using diamond slurry. Sintered samples were polished to 6 mm using diamond slurry and then polished to 1 mm electrochemical-mechanical polishing23  using  to reveal  grain  structure  and reduce SiC pullout. The microstructures were  characterized using scanning electron microcopy (SEM: Quanta,  FEI, Hillsborough, OR) along with energy-dispersive spectros copy (EDS: Pegasus 4000, EDAX, Mahwah, NJ) for elemental  analysis. Samples were also prepared for transmission electron  microscopy (TEM: Phillips CM200, FEI) using a focused ion  beam microscope (FIB: FEI Dual Beam 235, FEI). The crys tallographic analysis of the oxide was performed on an X-ray diffractometer with CuKa  radiation  (XRD: Rigaku  2500,  Tokyo, Japan). Densities of the oxide scale and phase analysis were measured from SEM micrographs at \\x02 1000 magnification  using  a  threshold and area measurement  routine  in Adobe  Photoshop (Fovea Pro 4, Reindeer Graphics).  III.  Results  (1)  Processing and Sample Characteristics  Samples were heated to 21001C and held for 5-9 min. allowing  at least 4 min past the last observed piston movement change to  ensure complete densification. Figure 1 is a plot of the relative  piston motion for the heating portion of the program for each of  the samples. The maximum piston motion was normalized so  that the final piston motion values were equal for ease of com parison. The first  sample to densify was HSWC followed by  HSWB and finally HS. HSWC reached its maximum piston motion at 20001C, while HSWB and HS did not reach full density until 21001C. The more rapid densification of  the samples  containing W is as  expected from previous data concerning containing WC additives.17,24 The  densification of diborides  densities of the HS, HSWC, and HSWB samples were measured to be 9.93, 10.14, and 10.30 g/cm3, while the theoretical values are 9.92, 9.99, and 9.91 g/cm3, respectively. The discrepancy be tween the measured densities and theoretical values for the W containing  samples  is due  to the  formation of high-density  phases  such as HfC and offset by some lower density oxides  as discussed below.  Figure  2  is  a  series of micrographs of  the  as-sintered HS  (Fig. 2(a)), HSWC (Fig. 2(b)), and HSWB (Figs. 2(c) and (d))  samples. The micrographs in Figs. 2(a)-(c) were obtained after  electrochemical polishing. The SiC is distinct and uniformly dis tributed throughout each of the samples. The SiC phase was de fined by the presence of only Si and C in the EDS spectrum and is  identified  as  the  darkest  phase  in  the  images  (labeled A in  Fig. 2(d)). The backscattered electron images revealed the grain  structure of  the HfB2 indicated by the presence of Hf-B in the EDS spectra (labeled B in Fig. 2(d)). The average grain size of  HfB2 in the HS, HSWC, and HSWB samples was 2.6, 1.8, and 1.7 mm, respectively. It was observed that the SiC morphology was  changed upon addition of WC or WB so that 20%-30% of the  SiC grains became rod-shaped with average aspect ratios of 2.9.  Table I.  Starting Powder Size and Purity  Powder  Supplier  Particle size  w  Purity (%)  Impurities (wt%)z  HfB2 SiC  Cerac (Milwaukee, WI)  \\x00325 mesh (44 mm)  99.5  Zr (0.2), Fe (0.02)  Alfa Aesar (Ward Hill, MA)  o1 mm  99.8  C (1.39), O (0.5),  N (0.13), Fe (0.0151),  Al (0.013), Mo (0.0134)  WC  Cerac  o1 mm  99.5  Cr (0.01), Fe (0.01)  WB  Cerac  \\x00325 mesh (44 mm)  99.5  Mg (0.08)  w  As received. zAs reported for any impurity over 0.01%.  Fig. 1. Plot of relative piston motion versus time for the HS, HSWC,  and HSWB samples. The onset of densification was more  rapid for  samples containing W. The vertical 21001C was attained.  line  indicates  the  time at which  August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2601  \\x0c', '2602  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  Fig. 2. SEM micrographs of the sintered (a) HS, (b) HSWC, and (c) HSWB samples. (d) Backscattered electron image of a region of HSWB highlighting  the five phases in the sample: (A) SiC, (B) HfB2, (C) W-C-O-B, (D) Hf-C-W, and (E) HfC.  Figure  2(d)  is  a backscattered electron SEM micrograph  evidence of  a WC or WB phase  remaining  in the  samples.  highlighting  the minor  phases  found  in  both HSWB and  At  least one additional phase is suggested by small peaks un HSWC. Grain pullout is evident in this sample as electrochem ical-mechanical polishing was not used to prepare the specimen.  Grains  containing W were  identified by EDS mapping con ducted with a  20 keV beam. The  grains highlighted in the  W map (labeled C in Fig. 2(d)) were subsequently analyzed by  point EDS analysis using a 10 keV beam. The 10 keV analyses of  at least 15 distinct areas revealed the presence of B, C, and O in  different concentrations with W. An additional phase composed  of Hf and C was identified through point EDS analysis. In some  instances (labeled D in Fig. 2(d)) the Hf-C phase contained W,  while in others (labeled E in Fig. 2(d)) no W was found. Image analysis of five EDS maps taken at \\x02 2000 combined with backscattered electron images revealed that on average 0.5% of the  sample was composed of the unidentified W-containing phase in  both the HSWC and HSWB samples. Because HfC has a higher density (12.2 g/cm3) than HfB2 (11.1 g/cm3) and WB (10.77 g/ cm3) its formation could contribute to the increase in the mea sured density of the HSWC and HSWB samples.  XRD patterns were taken of  the sintered samples  to deter mine the phases present (Fig. 3). Hexagonal HfB2 and cubic SiC are readily identified and are the only phases present in the HS  sample. In the HSWC and HSWB samples there is a shift in the  HfB2 peak attributable to a replacement of W on the Hf lattice. As expected, a replacement of the smaller W atom (1.41 A˚ ) for the Hf atom (1.58 A˚ ) on the hexagonal lattice shifts the peaks to higher 2y values with the (002) peak experiencing the greatest  shift of  all peaks within the observed angular  range. Peaks  matching the HfC pattern are found that are also shifted to higher 2y values  confirming the EDS findings of a Hf-W-C  phase. The peak shifts  are  shown in Table  II. There  is no  identifiable by any combination of W, C, O, Si, Hf, or B in the database. The weak peaks near 2y values of 301 and 381 may be  monoclinic HfO2 or peaks of the unknown phase(s). A ternary rhombohedral HfW4B5 or hexagonal (Hf,W)12B2\\x00x phase is possible, but not conclusively identified due to the low intensity  of the peaks and variability in the lattice parameters reported for these phases.25,26  Fig. 3. XRD of the sintered HS, HSWC, and HSWB samples showing  peaks of SiC (3C), HfB2, and HfC. The HfB2 and HfC peaks are both shifted to higher 2y values due to solid solutions formed with WC, WB,  or W. The peaks marked with a plus  symbol are unknown phase(s)  which may include monoclinic HfO2.  \\x0c', '(2)  Oxidation  Oxidation was  carried out at 16001, 18001, and 20001C in a  stagnant air atmosphere. The heating and cooling rates were 51C/min, so that the samples experienced oxidizing temperatures  before and after  the 30-min hold at maximum temperature.  Table III lists the average total oxide scale thickness as measured  from at least 10 points on each sample. The average total scale  thickness was measured by first performing EDS to determine  the boundary between oxidized grains and the unaltered bulk.  A measurement was then taken using the ruler tool  in the xT  microscope control software of the SEM. The oxide thickness values at 16001 and 18001C are within one standard deviation of  each other for all samples suggesting no improvement or detri ment to oxidation resistance by adding W-containing phases at these temperatures. While at 20001C both HSWC and HSWB had oxide scale thicknesses that were 430% thinner than the  HS sample. As a comparison, samples of HfB2-20 vol% SiC-3 vol% WC heated to 16001 and 18001C had average total scale thicknesses of 47 and 73 mm, respectively.  (A)  Oxidation at 16001C:  HSWC and HSWB samples heated to 16001C produced oxide scales with the same morphol ogy and chemistry of the scale observed on the HS samples heated to 16001C. The oxide scales were composed of an outer  layer of SiO2 and a porous HfO2 layer underneath. EDS analysis of the SiO2 scale in each of the samples revealed Al impurities whose concentration varied throughout the scale, but was never 41 mol% assuming Al as  the only impurity in SiO2. W-containing phases were observed as submicrometer spherical grains  scattered throughout the HfO2 layer in the oxidized HSWC and HSWB samples. Additionally, a W-containing phase was found  between the HfO2 grains grain pullout due to polishing allowed its observation. Point  (Fig. 4)  in areas where porosity or  EDS analysis of the features is limited by their size, but exam ination of at  least 10 regions with the  same  features always  showed the presence of W and O and sometimes Hf. The Hf  signal may be from the background, while any potential small Si  (K, 1.74 kV) peak may be obscured by the Hf (M, 1.645 kV) and  W (M, 1.775 kV) peaks that bookend it  in the energy spectra.  Monoclinic HfO2 and hafnon (HfSiO4) analysis for each of the samples. HfSiO4 is the reaction product from the combination of HfO2 and SiO2. Calculated HfO2- SiO2 phase diagrams show that HfSiO4 is stable up to B17261C.27,28  are  found by XRD  (B)  Oxidation at 18001C:  Representative images of the oxide scales formed on the samples after exposure to 18001C are  found in Fig. 5. Point EDS analysis showed four layers that are  labeled in Fig. 5:  (I) dense SiO2 outer HfO2 grains, (II) porous HfO2 penetrated by SiO2, (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C, and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C. The average length of each layer was calculated  layer with distributed  using the same method as the total scale length to find that layers (I)-(III) of all samples were all within 3 mm. The porosity of  all the layers was also similar; in particular, culated to be 90%71% dense  layer (III) was cal for all  samples.  In measuring  porosity by the  thresholding method,  the  inside of  the pores  were taken as part of the matrix, so the actual density would be  somewhat lower, but the same method was used for all samples.  Layer (IV) is commonly known as the depleted layer and has been reported for ZrB2-SiC and HfB2-SiC systems.8,13-17,21 The average thickness of the depleted layer in the HS samples was 6 mm; however, samples HSWC and HSWB did not have a layer  (IV); the entire porous interface between the bulk and the oxide is comprised of HfO2. As in the 16001C oxidized samples, SiO2 scale in each of the samples had Al impurities whose concentration varied throughout the scale, but was never 41 mol%  the  assuming Al as the only impurity in SiO2. W-containing grains are present throughout all the layers in  both HSWC and HSWB. Examples of the W-containing grains are indicated by arrows in Fig. 5(d). EDS of the larger (1-2 mm)  grains  showed the presence of W and O. A third phase com prised of Hf, Si, and O is recognizable between HfO2 grains and the amorphous SiO2. The phase is likely HfSiO2 (labeled in Fig. 5(d)) as evidenced by its presence at 16001C and the increase  in the peak intensity observed in the XRD patterns of all the samples after oxidation to 18001C. Besides HfSiO4, only monoclinic and tetragonal HfO2 are present in the XRD patterns of all samples heated to 18001C.  (C)  Oxidation at 20001C:  Figure 6 is a series of com bined micrographs taken of the oxide scale resulting from exposure to 20001C. The layers of oxide scale for the HS sample heated to 20001C are: (I) a dense SiO2 outer layer with distributed HfO2 grains; (II) porous HfO2 penetrated by SiO2 (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C; and (IV) porous HfB2 with inclusions of different concentrations of Si, O, and C. The total scale thickness (Table I)  Table II.  Peak Shifts Observed for HfB2 and HfC in Samples HSWC and HSWB  Peak  Peak shift in 2y  HSWC  HSWB  HfB2 (001) HfB2 (100) HfB2 (101) HfB2 (002) HfB2 (110) HfC (111)  0.01  0.04  0.01  0.03  0.03  0.06  0.05  0.09  0.04  0.05  0.5  0.3  HfC (200)  0.5  0.3  HfC (220)  0.7  0.5  Table III.  Oxide Scale Thickness  Temperature (1C)  Oxide scale thickness (mm)  HS  HSWC  HSWB  1600  3573.2 7875.5 826756.1  3775.3 6775.5 537728.1  3472.3 7677.2 565723.8  1800  2000  Fig. 4. SEM micrograph of the W-containing phases that are found in the oxide scale after samples HSWC and HSWB are oxidized to 16001C.  Black arrows indicate the intergranular phase found between HfO2 that is composed of W-O and possibly Hf, while the white arrow indicates  spherical grains that are found throughout the oxide scale that also con tain W-O and possibly Hf.  August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2603  \\x0c', '2604  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  Fig. 5. SEM micrographs of the (a) HS, (b) HSWC, and (c) HSWB samples after oxidation at 18001C. The white arrows indicate the location of W containing grains in HSWC (d) while the HfSiO4, HfO2, and SiO2 phases are labeled. HfSiO4 was found in all three samples oxidized to 18001C. The oxide scale layers are indicated in (a) as (I) dense SiO2 outer layer with distributed HfO2 grains, (II) porous HfO2 penetrated by SiO2, (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C, and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C.  was  reduced by the addition of WC and WB while the SiO2 penetration into the scale and the density increased. In the  HSWC and HSWB samples  the distinction between layer  (I)  and layer  (II)  is blurred through the  extensive  formation of  HfSiO4. Layer (III) is reduced from an average thickness of 357 mm in HS to 170 mm in HSWC and 93 mm in HSWB. In the  Fig. 6. SEM micrographs of the (a) HSWB, (b) HSWC, and (c) HS samples after oxidation at 20001C. The layers: (I) a dense SiO2 outer layer with distributed HfO2 grains; (II) porous HfO2 penetrated by SiO2 and HfSiO4; (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C; and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C are indicated in the images. The magnified image in (a) shows the SiO2 and HfSiO4/HfO2 in layer (I) and (II), while the magnified image in (b) shows W-rich phases found in layer (III) as indicated by the arrows.  \\x0c', 'August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2605  Fig. 7.  (a) SEM micrograph showing three distinct glass phases of the HSWB sample oxidized to 20001C: (A) W-rich, (B) SiO2 with Al, Fe, Ca, and Ba, (C) SiO2-rich, and (D) HfO2. (b) A TEM section from layer (I). The dark phase is W-rich. (c) TEM micrograph of a section that was prepared after the surface glass had been removed. The inset is the SEM micrograph of the surface of the sample showing the Pt cap applied for FIB cutting. The HfO2 (the light phase in the SEM and the dark phase in TEM) was monoclinic as determined by selected area electron diffraction (not shown). (d) TEM-EDS of  the glass phase showing the composition of the light and dark glasses in (c).  HSWC (Fig. 6(b)) and HSWB (Fig. 6(a)) the average thickness of the depleted later (layer (IV)) is 22 mm in HSWC and 32 mm in HSWB compared with 90 mm in HS. Image analysis calculated a  layer (III) density of 87%, 96%, and 97% in HS, HSWC, and  HSWB, respectively.  The SiO2 in layer (I) of the HS sample had impurities of Al, Ca, Ba, and Fe whose concentration varied spatially. Addition ally, the SiO2 in the oxide scale contains the same impurities (Al, Ca, Fe, Ba) as were found in the HS sample but were also shown  to form phases with W. Figure 6(a) is a magnified view of layers  (I) and (II). Extensive HfSiO4 formation is observed surrounding the lighter HfO2 grains. The HfSiO4 does not contain W, but small pockets (indicated by an arrow) of Hf-Si-W-O can be  found in layer  (II) of both the HSWC and HSWB samples.  Figure 6(b) is a magnified image of layer (III) showing the W containing spheres that were found scattered throughout in both  the HSWC and HSWB samples.  SEM and TEM images of  the W-containing phases  in the  glass are shown in Fig. 7. In Fig. 7(a) SEM-EDS suggests three  distinct glass phases:  (2) SiO2 with Al, Fe, Ca, and Ba (labeled B), and (3) SiO2-rich (labeled C). A TEM section from layer (I) shows a similar microstructure for  (1) W-rich (labeled A),  taken of the glass phase. Fitting the EDS pattern under the as sumption that all the impurities are oxides, the light glass contains o0.6 mol% of any impurity with HfO2 having the highest concentration. The dark glass contained Si, O, Hf, W, Ca, Al, Ba, Fe, and possibly Mg. The low intensity of the MgKa peak  and its overlap with an Hf M peak make Mg difficult to discern  in this  spectra. Selected area electron diffraction analysis  re vealed all phases in Fig. 7(b) and those in 7(c), exclusive of the  HfO2, to be amorphous. The XRD patterns of the sample surfaces after oxidation at 20001C reveal prominent HfSiO4 peaks in the HSWB sample, which correlates well with the observed microstructure. Approximately 60% of the top surface of the 20001C oxidized HSWB is  comprised of  the dense HfSiO4/Hf-Si-W-O phase while localized regions of dense HfSiO4/HfO2 comprise approximately 30% of the top surface observed in 20001C oxidized HSWC.  Although HfO2 would be expected to transform to tetragonal at 20001C, the slow cooling rate (51C/min) would allow conversion  to the monoclinic phase at lower temperatures and account for the small tetragonal peak (o15% of the height of the monoclinic (\\x00111) peak) observed for all the samples.  the multiphase glass with the black phase being W-rich. Figure  7(c) is an overview of a section that was prepared after the sur face glass had been removed. The inset is the SEM micrograph  IV.  Discussion  of the surface of the sample showing the Pt cap applied for FIB  After  sintering, no pure WB or WC phases were observed  cutting. The HfO2 phase in TEM) was monoclinic as determined by selected area  (the light phase in the SEM and the dark  electron diffraction (not  shown). TEM-EDS (Fig. 7(d)) were  in the HSWC and HSWB samples. The W was  incorporated  in the HfB2 and HfC matrix as evidenced by their peak shifts in the XRD spectra. HfC can be formed by a reaction of HfO2  \\x0c', 'and excess C. HfO2 is present as oxygen contamination on the HfB2 powders and oxidation during milling. Amorphous borides  (B2O3) also exist at from wear of the high-density  the surface. C can be introduced  polyethylene  bottles  during  milling or  from the graphite dies used to sinter  the powders.  The oxides may be consumed at elevated temperatures through  a reaction with free carbon by reactions (3) or (4) or with WC  (reaction [5]).  HfO2 ðsÞ þ B2O3 ðl Þ þ 5CðsÞ ! HfB2 ðsÞ þ 5COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x008:6 KJ=mol  (3)  HfO2 ðsÞ þ 3CðsÞ ! HfCðsÞ þ 2COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x00111:5 KJ=mol  (4)  HfO2 ðsÞ þ 3WCðsÞ ! HfCðsÞ þ 3WðsÞ þ 2COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x00489:3 KJ=mol  (5)  The presence of HfC after sintering diborides has been shown to be dependent on sintering temperature and additives.29,30  The phase diagram of HfB2 and WBB2 suggests a maximum 23 mol% at B23651C and W concentration of 10 mol% or above at temperatures above 12001C,31 while melting exper iments suggest  that approximately 4 mol% W can be incorpolattice.32 The same crystalline lattice and  rated into the HfB2 similar size allow 40 mol% of WC to be dissolved in HfC at 20271C.33 Limited literature exists for the HfC-WB system, but  the similar XRD peak shift values of HSWC and HSWB suggest  solubility of W in HfC in the HSWB sample. The W that  is  not incorporated into either HfC or HfB2 can be found in grains shown to contain W, O, C, and B. The addition of W-containing  phases  increases  the  sinterability  of  the  samples,  results  in smaller HfB2 SiC grains. A similar SiC microstructure was shown by Zhang et al.34 for pressureless sintered ZrB2-SiC with B4C additions and WC impurities. However, Chamberlain et al.24 do not show  grains,  and promotes  evolution of  acicular  the same SiC microstructure for WC impurities in hot pressed  ZrB2-SiC. It is possible that excess B or C (from WB, WC, or B4C34) in the system can promote the b (cubic) to a (hexagonal) transformation at temperatures below 18001C, with the a-SiC having acicular grain morphology.35 Unfortunately the SiC XRD peak has a very low intensity and the main b and a-SiC peaks are within 0.071, so phase identification is imprac tical. Engineering the microstructure may prove beneficial as it  has been shown that the size and morphology of HfB2 grains can affect strength.29,35,36 The oxidation of HS, HSWC, and HSWB samples resulted  the SiC and  in oxide scales with comparable total thicknesses and scale morphologies at 16001 and 18001C. The difference in the aver age thickness of the total oxide scale for each sample was within  the spread measured for an individual sample. The oxide scales formed on the HSWC and HSWB samples at 20001C were sig nificantly different  than the scale formed on HS. The HSWC  and HSWB samples had less porous scales that were 30% thin ner than the scale formed on the HS sample. The three methods  by which W-additions may impact oxidation resistance are (1)  controlling the phase transformations of the HfO2 scale, (2) increasing the viscosity of the SiO2 scale, or (3) promoting densification of the HfO2 scale. The addition of W neither promoted nor delayed the monoclinic to tetragonal phase transformation,  thus not affecting the density of the HfO2 by this method. Because no W was found in the SiO2 at 16001 and 18001C by SEM-EDS the addition of W-containing phases would not be  expected to alter  the viscosity or melting temperature of  the  SiO2 glass and the SiO2-rich layer would be expected to be protective at 16001 and 18001C.3,4,18,21 A two-phase glass is not 18001C which (17231C),  observed  until  above  is  above  the melting  temperature of pure SiO2 form as the glass cools. Borate and silicate glasses with Group  so the  two phase may  IV-VI transition metal (Hf, Zr, Ti, Ta, W) oxides tend to exhibit microphase separation and crystallization upon cooling.15,37,38  Glass  compositions  that  exhibit phase  separation posses  an  increased viscosity in the single-phase liquid.38  two-phase  region compared with a  Only a few phase diagrams for W and Hf in B2O3 or SiO2 are reported. The most studied diagram, HfO2-SiO2, reveals that at 20001C a single liquid phase exists for compositions with  approximately 60 mol% (or greater) of SiO2 and that at lower SiO2 concentrations the Si-Hf-O liquid phase is found in equilibrium with HfO2.27,37 In the WO3-B2O3 system a single-phase liquid is formed at all compositions above 14301C.39 No phase  diagram exists for HfO2-B2O3 but a study of glass melts found that o1 mol% HfO2 could be dissolved in B2O3.40 However, the same study found that sodium borosilicate or aluminum  borosilicate glasses can dissolve up to 17 mol% HfO2 with the allowable HfO2 concentration dependent on the concentration of the other glass species.39 As evidenced by Fig. 7 the glass chemistry found in the 20001C sample has spatial variances in  chemistry. The presence of Al, Na, Ca, Ba, and Fe in the glass  could originate from the impurities in the starting powders or  be incorporated from the Ca-stabilized ZrO2 crucible that holds the sample. Group I and Group II impurities have been  observed in our previous experiments and also reported in the literature,16 but have not received much attention to date. How ever,  as  the development of new UHTC compositions  seek  to add elements  in order  to improve glass properties  such as  viscosity the presence of these impurities may play a crucial role  in the outcome of these efforts.  The  phase  diagrams  for WO3 with SiO2 temperatures of interest, but  or  B2O3 inspection of  are  unavailable at  the  other  transition metal oxides with SiO2 such as Nb2O5-SiO2 and Cr2O3-SiO2 suggest two-phase glasses above 16951 and 22001C, respectively.41,42 Therefore the observed existence of a  two-phase  region  in  the WO3-SiO2 glass may improve oxidation resistance by the glass at 20001C and/or during  system seems  logical.  The  two-phase  increasing the viscosity of  its slow cool down.  The  impact of densification in the HfO2-based layers was examined in the HSWC and HSWB samples. The HfO2 contained W both as individual grains and as a phase between  HfO2 grains. Reactions (6)-(8) describe the possible oxidation of WC and WB.43,44  WCðcrÞ þ 5=2 O2 ðgÞ ! WO3 ðcrÞ þ CO2 ðgÞ  (6)  2 WBðcrÞ þ 9=2 O2 ðgÞ ! 2 WO3 ðcrÞ þ B2O3 ðl Þ  (7)  WO3 ðcrÞ ! WO3 ðgÞ  (8)  The HfO2-WO3 phase diagram shows a solubility of WO3 of about 5 mol% at 16001-20001C, while above 5 mol% the HfO2 solid solution will be accompanied by a liquid.22 The solubility  of W in HfB2 and HfC is between 4 and 40 mol% suggesting localized regions of W concentrations necessary for phase sep aration could exist. The resulting liquid has a melting point of 14731C. The presence of  this  liquid can promote liquid-phase  sintering of  the HfO2 scales for the HSWC and HSWB samples. The reduction in the length of the total SiC-depleted layer at 18001C in the samples  resulting in the observed denser oxide  containing W can be explained by a reduction in oxygen pen etration through the denser oxide  scale. Upon cooling,  the  resulting crystalline phases are a HfO2 solid solution with o5 mol% WO3 and WO3 with no solid solubility for HfO2. SEM-EDS confirmed the existence of W-rich phases between  the HfO2 grains in the postoxidation analysis. In addition to the denser HfO2 layer (layer (III)), layer (II) in the samples with W additions is composed of HfSiO4 and HfO2. Although a dense HfSiO4 would be expected to have a lower oxygen diffusion coefficient than the porous SiO2-filled HfO2 regions of HS, the decomposition  found  in  the  outermost  2606  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  \\x0c', 'temperature of HfSiO4 is 17261C27 so that the HfSiO4 is likely formed on cooling and would not be responsible for the decrease in oxygen penetration at 20001C.  The  combined effect of  the more viscous outer  layer and  denser inner layer promote oxidation resistance and provide for an overall thinner scale at exposure temperatures of 20001C in  HSWC and HSWB.  V.  Conclusion  Dense HfB2-SiC samples were prepared with additives of WC and WB. Both the WC and WB additives promoted sintering  while reducing the grain size of HfB2. Solid solutions of W-Hf- B and W-Hf-C were formed in both samples. The WCand WB-containing HfB2-SiC samples oxidized to 16001 and 18001C exhibited similar oxide scales as the HfB2-SiC sample. However, the samples with WC and WB showed a 30% reduction in scale thickness when the samples were oxidized at 20001C  due to a more viscous phase separated glass found in the out ermost  regions  of  the  scale  and  a  denser  inner HfO2  that  restricted oxygen penetration to the sample.  Acknowledgment  The authors would like to thank Pavel Mogilevsky for his help in preparing the  TEM-FIB sample and useful discussions regarding TEM-EDS.  References  1R. L. Cloughtery, R. L. Pober, and L. Kaufman, ‘‘Thermogravimetric Study of the Oxidation of ZrB2 in the Temperature Range of 800-15001C,’’ Trans. Met. 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  "PDF": "Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC.pdf",
  "Text": "['Oxidation Resistance of Hafnium Diboride Ceramics with Additions of  Silicon Carbide and Tungsten Boride or Tungsten Carbide  Carmen M. Carney,*,w,z,y  Triplicane A. Parthasarathy,*,z,y  and Michael K. Cinibulk*,z  zAir Force Research Laboratory, Wright-Patterson Air Force Base, Ohio 45433-7817  y  UES Inc., Dayton, Ohio 45432  Dense samples of HfB2-SiC, HfB2-SiC-WC, and HfB2-SiC- WB were prepared by field-assisted sintering. The WB and WC  additives were incorporated by solid solution into the HfB2 and the HfC that formed during sintering. Oxidation of the samples furnace oxidation between 16001 was studied using isothermal and 20001C. Sample microstructure and chemistry before and after oxidation were analyzed by scanning electron microscopy  and X-ray diffraction. The addition of WC and WB did not alter  oxidation kinetics of the baseline HfB2-SiC composition below 18001C; however, at 20001C, HfB2-SiC-WC and HfB2-SiC- WB had oxide scales that were 30% thinner than the oxide  scale of HfB2-SiC. It liquid-phase densification of  is believed that WC and WB promoted the HfO2 the path of oxygen ingress, during oxidation.  scale,  thereby reducing  I.  Introduction  TRANSITION metal borides and carbides with melting temperatures exceeding 27001C are commonly referred to as ultra-high-temperature ceramics (UHTCs). Of these materials,  hafnium diboride (HfB2) has shown the highest oxidation resistance.1-5 Although its density is higher than that of zirconium diboride (ZrB2)1-4 (10.5 vs 6.09 g/cm3) its high melting point (431001C) and high thermal conductivity (B100 W/MK  at room temperature) make it a leading candidate for extreme  environment thermal protection systems such as those found at the sharp leading edges of hypersonic vehicles.6,7  The oxidation of HfB2 proceeds according to reaction (1) where the B2O3 liquid formed is volatile and expected to be removed through evaporation above 11001C. Although HfO2 is a stable refractory with a melting point of 27501C the oxide  formed on the surface of the diboride is porous and nonprotective from further oxygen penetration.3 As  such, development  of these materials in the 1960s led to the addition of 5-30 vol%  of silicon carbide (SiC) to form a borosilicate glass upon oxidation.1,3-5 Above 11001C, SiC will oxidize by reaction (2).  Available B2O3 SiO2 borosilicate glass that covers or fills the pores of the HfO2 scale. The SiO2 glass has shown to improve oxidation resistance (o17001C) when at moderate temperatures the oxidation of pure HfB2.3,6,7 HfB2 ðcrÞ þ 5=2 O2 ðgÞ ! HfO2 ðcrÞ þ B2O3 ðl Þ  can  dissolve  in  to  form a  protective  compared with  (1)  SiCðcrÞ þ 3=2 O2 ðgÞ ! SiO2 þ COðgÞ  (2)  However, as these materials have been tested in environments  that more closely simulate atmospheric reentry and hypersonic cruise applications that include higher temperatures (417231C) and flowing air8-13 it has become apparent that there is a point  at which the protective SiO2-based scale begins to fail. Above the melting point of SiO2 (17231C) it is observed that the less viscous SiO2 will flow into the pores of the HfO2 and even flow from the sample resulting in a less protective porous outer scale.12 Recent research has focused on extending the oxidation  resistance of HfB2-SiC to these temperatures by adding additional phases.7,13-17 There are three strategies for selecting the  appropriate additive that will accomplish one or more goals of  (1) controlling the phase transformations of the HfO2 scale, (2) increasing the viscosity of the SiO2 scale, or (3) promoting densification of the HfO2 scale. Mixed results have been obtained with the addition of rare earth elements (Ta, La and Nd) and some transition metals (Ti,  Cr, and Cr) with reduced oxidation being observed for oxidation temperatures below 17001C7,15-21 and increased oxidation above 17001C.12,15 Most recently, Zhang et al.17,21 have prepared ZrB2based UHTCs with additions of WC to promote liquid-phase  sintering in the oxide scale. It was shown that  the sample con taining WC had a denser ZrO2 scale and the weight gain during oxidation was reduced at temperatures of 16001C and below.  The present paper focuses on the addition of tungsten (W) to  an HfB2-based UHTC to promote improved oxidation resistance at temperatures 416001C. The WO3-HfO2 phase diagram suggests a solid solution between the phases of at least 5 mol% temperatures up to and beyond 20001C.22 Similar  WO3 at ZrO2-WO3, HfO2 will form a liquid phase with WO3 at temperatures above 12801C.22 In the present study, W was added in  to  two phases:  (1) WC and (2) WB. Samples with W-based addi tives are  compared with a baseline  composition of HfB2-15 vol% SiC. SiC content was held constant in the samples to allow  a direct  comparison. Fifteen volume percent SiC was  chosen  based on internal experiments that showed it to provide oxida tion scales of similar total  lengths and comparable weight gains  as 20 vol% SiC at  the tested temperatures, but with a thinner 16001  SiO2 scale. The samples were 20001C. Because of  oxidized  between  and  the differences  in weight between the for mation of HfO2 and WO3 volatility of WO3, thickness was used as a measure of oxidation resistance.  and the  scale  II.  Experimental Procedure  (1)  Powder Processing  Commercially available HfB2, b-SiC, WC, and WB powders were used to prepare the samples. Table I is a list of the sup pliers, purity, and starting particle sizes. Powders were used as  received except  for the HfB2 which was premilled using Si3N4 grinding media in isopropanol for 60 h to reduce the grain size.  The weight change of the grinding media was 0.07 g, which was  0.01 wt% of  the total batch weight. Representative batches of  milled HfB2 were measured using a laser diffraction particle size  N. Jacobson—contributing editor  This work was financially supported under Contract No. FA8650-10-D-5226.  *Member, The American Ceramic Society.  w  Author  to whom correspondence  should  be  addressed.  e-mail:  carmen.carney@  wpafb.af.mil  Manuscript No. 28917. Received November 19, 2010; approved January 21, 2011.  Journal  J. Am. Ceram. Soc., 94 [8] 2600 - 2607 (2011)  DOI: 10.1111/j.1551-2916.2011.04462.x  r 2011 The American Ceramic Society  2600  \\x0c', 'analyzer (LS230, Beckman Coulter, Brea, CA) and had an average particle size of 1.3 mm (D90 5 2.2 mm). The oxygen content of the premilled HfB2 was 0.86 wt% compared with 0.58 wt% in the as received powder as measured by Leco Corporation (St.  Joseph, MI). Compositions were chosen based on the HfO2- WO3 phase diagram. The prepared samples were HfB2-15 vol% SiC (HS), HfB2-15 vol% SiC-3 vol% WC (HSWC) and HfB2- 15 vol% SiC-3 vol% WB (HSWB). The powder mixtures were  ball milled in isopropanol  for 24 h with SiC grinding media,  dried at  room temperature, and subsequently dry milled for  12 h. Typical weight loss of the SiC grinding media after milling  the additives with HfB2 was 0.2 mg (0.1 wt% of the total batch). The powders were sieved through an 80-mesh (177 mm) screen.  (2)  Sintering and Sample Preparation  The milled powders were loaded into a 25-mm-diameter graphite  die. A layer of BN and graphite foil separated the powder from  the die with the powder  in contact with the graphite foil. The  powder-filled dies were cold pressed at approximately 20 MPa  and loaded into the field-assisted sintering (FCT Systeme GmbH,  Model HPD 25-1, Rauenstein, Germany) unit. The graphite die  was wrapped in graphite felt to limit heat radiation from the die. Heating and cooling rates were 501C/min. and 32 MPa was applied while heating up to 16001C. The 32 MPa was held for the  remainder of the sintering schedule and rereleased during the free cool (below 10001C). The samples were held at 21001C for times  between 5 and 9 min. The temperature was measured by a py rometer focused on the bottom of a bore hole in the upper punch B5 mm from the surface of the powder. Densification was mon itored by tracking the movement of the pistons.  Sintered samples were cut with a wire electron discharge 2.5 mm \\x02 2.0 mm \\x02 9.0 mm rectangles machine into and polished using diamond slurry to a 1 mm finish on all six sides  using an autopolisher to ensure samples with parallel sides and  uniform sizes.  (3)  Oxidation Exposure and Sample Characterization  Polished samples were heated in a zirconia element furnace (ZrF-25: Shinagawa Refractories Co., Tokyo, Japan) to 16001, 18001, or 20001C and held at temperature for 30 min. The heating and cooling rates were 51C/min., which were limited by the  furnace. Samples were placed on a concave ZrO2 crucible to limit contact between the sample and ZrO2. Oxidized samples were polished in cross section perpendicular to the bottom (2.5 mm \\x02 9.0 mm side facing the crucible) of the sample to a 1 mm finish using diamond slurry. Sintered samples were polished to 6 mm using diamond slurry and then polished to 1 mm electrochemical-mechanical polishing23  using  to reveal  grain  structure  and reduce SiC pullout. The microstructures were  characterized using scanning electron microcopy (SEM: Quanta,  FEI, Hillsborough, OR) along with energy-dispersive spectros copy (EDS: Pegasus 4000, EDAX, Mahwah, NJ) for elemental  analysis. Samples were also prepared for transmission electron  microscopy (TEM: Phillips CM200, FEI) using a focused ion  beam microscope (FIB: FEI Dual Beam 235, FEI). The crys tallographic analysis of the oxide was performed on an X-ray diffractometer with CuKa  radiation  (XRD: Rigaku  2500,  Tokyo, Japan). Densities of the oxide scale and phase analysis were measured from SEM micrographs at \\x02 1000 magnification  using  a  threshold and area measurement  routine  in Adobe  Photoshop (Fovea Pro 4, Reindeer Graphics).  III.  Results  (1)  Processing and Sample Characteristics  Samples were heated to 21001C and held for 5-9 min. allowing  at least 4 min past the last observed piston movement change to  ensure complete densification. Figure 1 is a plot of the relative  piston motion for the heating portion of the program for each of  the samples. The maximum piston motion was normalized so  that the final piston motion values were equal for ease of com parison. The first  sample to densify was HSWC followed by  HSWB and finally HS. HSWC reached its maximum piston motion at 20001C, while HSWB and HS did not reach full density until 21001C. The more rapid densification of  the samples  containing W is as  expected from previous data concerning containing WC additives.17,24 The  densification of diborides  densities of the HS, HSWC, and HSWB samples were measured to be 9.93, 10.14, and 10.30 g/cm3, while the theoretical values are 9.92, 9.99, and 9.91 g/cm3, respectively. The discrepancy be tween the measured densities and theoretical values for the W containing  samples  is due  to the  formation of high-density  phases  such as HfC and offset by some lower density oxides  as discussed below.  Figure  2  is  a  series of micrographs of  the  as-sintered HS  (Fig. 2(a)), HSWC (Fig. 2(b)), and HSWB (Figs. 2(c) and (d))  samples. The micrographs in Figs. 2(a)-(c) were obtained after  electrochemical polishing. The SiC is distinct and uniformly dis tributed throughout each of the samples. The SiC phase was de fined by the presence of only Si and C in the EDS spectrum and is  identified  as  the  darkest  phase  in  the  images  (labeled A in  Fig. 2(d)). The backscattered electron images revealed the grain  structure of  the HfB2 indicated by the presence of Hf-B in the EDS spectra (labeled B in Fig. 2(d)). The average grain size of  HfB2 in the HS, HSWC, and HSWB samples was 2.6, 1.8, and 1.7 mm, respectively. It was observed that the SiC morphology was  changed upon addition of WC or WB so that 20%-30% of the  SiC grains became rod-shaped with average aspect ratios of 2.9.  Table I.  Starting Powder Size and Purity  Powder  Supplier  Particle size  w  Purity (%)  Impurities (wt%)z  HfB2 SiC  Cerac (Milwaukee, WI)  \\x00325 mesh (44 mm)  99.5  Zr (0.2), Fe (0.02)  Alfa Aesar (Ward Hill, MA)  o1 mm  99.8  C (1.39), O (0.5),  N (0.13), Fe (0.0151),  Al (0.013), Mo (0.0134)  WC  Cerac  o1 mm  99.5  Cr (0.01), Fe (0.01)  WB  Cerac  \\x00325 mesh (44 mm)  99.5  Mg (0.08)  w  As received. zAs reported for any impurity over 0.01%.  Fig. 1. Plot of relative piston motion versus time for the HS, HSWC,  and HSWB samples. The onset of densification was more  rapid for  samples containing W. The vertical 21001C was attained.  line  indicates  the  time at which  August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2601  \\x0c', '2602  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  Fig. 2. SEM micrographs of the sintered (a) HS, (b) HSWC, and (c) HSWB samples. (d) Backscattered electron image of a region of HSWB highlighting  the five phases in the sample: (A) SiC, (B) HfB2, (C) W-C-O-B, (D) Hf-C-W, and (E) HfC.  Figure  2(d)  is  a backscattered electron SEM micrograph  evidence of  a WC or WB phase  remaining  in the  samples.  highlighting  the minor  phases  found  in  both HSWB and  At  least one additional phase is suggested by small peaks un HSWC. Grain pullout is evident in this sample as electrochem ical-mechanical polishing was not used to prepare the specimen.  Grains  containing W were  identified by EDS mapping con ducted with a  20 keV beam. The  grains highlighted in the  W map (labeled C in Fig. 2(d)) were subsequently analyzed by  point EDS analysis using a 10 keV beam. The 10 keV analyses of  at least 15 distinct areas revealed the presence of B, C, and O in  different concentrations with W. An additional phase composed  of Hf and C was identified through point EDS analysis. In some  instances (labeled D in Fig. 2(d)) the Hf-C phase contained W,  while in others (labeled E in Fig. 2(d)) no W was found. Image analysis of five EDS maps taken at \\x02 2000 combined with backscattered electron images revealed that on average 0.5% of the  sample was composed of the unidentified W-containing phase in  both the HSWC and HSWB samples. Because HfC has a higher density (12.2 g/cm3) than HfB2 (11.1 g/cm3) and WB (10.77 g/ cm3) its formation could contribute to the increase in the mea sured density of the HSWC and HSWB samples.  XRD patterns were taken of  the sintered samples  to deter mine the phases present (Fig. 3). Hexagonal HfB2 and cubic SiC are readily identified and are the only phases present in the HS  sample. In the HSWC and HSWB samples there is a shift in the  HfB2 peak attributable to a replacement of W on the Hf lattice. As expected, a replacement of the smaller W atom (1.41 A˚ ) for the Hf atom (1.58 A˚ ) on the hexagonal lattice shifts the peaks to higher 2y values with the (002) peak experiencing the greatest  shift of  all peaks within the observed angular  range. Peaks  matching the HfC pattern are found that are also shifted to higher 2y values  confirming the EDS findings of a Hf-W-C  phase. The peak shifts  are  shown in Table  II. There  is no  identifiable by any combination of W, C, O, Si, Hf, or B in the database. The weak peaks near 2y values of 301 and 381 may be  monoclinic HfO2 or peaks of the unknown phase(s). A ternary rhombohedral HfW4B5 or hexagonal (Hf,W)12B2\\x00x phase is possible, but not conclusively identified due to the low intensity  of the peaks and variability in the lattice parameters reported for these phases.25,26  Fig. 3. XRD of the sintered HS, HSWC, and HSWB samples showing  peaks of SiC (3C), HfB2, and HfC. The HfB2 and HfC peaks are both shifted to higher 2y values due to solid solutions formed with WC, WB,  or W. The peaks marked with a plus  symbol are unknown phase(s)  which may include monoclinic HfO2.  \\x0c', '(2)  Oxidation  Oxidation was  carried out at 16001, 18001, and 20001C in a  stagnant air atmosphere. The heating and cooling rates were 51C/min, so that the samples experienced oxidizing temperatures  before and after  the 30-min hold at maximum temperature.  Table III lists the average total oxide scale thickness as measured  from at least 10 points on each sample. The average total scale  thickness was measured by first performing EDS to determine  the boundary between oxidized grains and the unaltered bulk.  A measurement was then taken using the ruler tool  in the xT  microscope control software of the SEM. The oxide thickness values at 16001 and 18001C are within one standard deviation of  each other for all samples suggesting no improvement or detri ment to oxidation resistance by adding W-containing phases at these temperatures. While at 20001C both HSWC and HSWB had oxide scale thicknesses that were 430% thinner than the  HS sample. As a comparison, samples of HfB2-20 vol% SiC-3 vol% WC heated to 16001 and 18001C had average total scale thicknesses of 47 and 73 mm, respectively.  (A)  Oxidation at 16001C:  HSWC and HSWB samples heated to 16001C produced oxide scales with the same morphol ogy and chemistry of the scale observed on the HS samples heated to 16001C. The oxide scales were composed of an outer  layer of SiO2 and a porous HfO2 layer underneath. EDS analysis of the SiO2 scale in each of the samples revealed Al impurities whose concentration varied throughout the scale, but was never 41 mol% assuming Al as  the only impurity in SiO2. W-containing phases were observed as submicrometer spherical grains  scattered throughout the HfO2 layer in the oxidized HSWC and HSWB samples. Additionally, a W-containing phase was found  between the HfO2 grains grain pullout due to polishing allowed its observation. Point  (Fig. 4)  in areas where porosity or  EDS analysis of the features is limited by their size, but exam ination of at  least 10 regions with the  same  features always  showed the presence of W and O and sometimes Hf. The Hf  signal may be from the background, while any potential small Si  (K, 1.74 kV) peak may be obscured by the Hf (M, 1.645 kV) and  W (M, 1.775 kV) peaks that bookend it  in the energy spectra.  Monoclinic HfO2 and hafnon (HfSiO4) analysis for each of the samples. HfSiO4 is the reaction product from the combination of HfO2 and SiO2. Calculated HfO2- SiO2 phase diagrams show that HfSiO4 is stable up to B17261C.27,28  are  found by XRD  (B)  Oxidation at 18001C:  Representative images of the oxide scales formed on the samples after exposure to 18001C are  found in Fig. 5. Point EDS analysis showed four layers that are  labeled in Fig. 5:  (I) dense SiO2 outer HfO2 grains, (II) porous HfO2 penetrated by SiO2, (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C, and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C. The average length of each layer was calculated  layer with distributed  using the same method as the total scale length to find that layers (I)-(III) of all samples were all within 3 mm. The porosity of  all the layers was also similar; in particular, culated to be 90%71% dense  layer (III) was cal for all  samples.  In measuring  porosity by the  thresholding method,  the  inside of  the pores  were taken as part of the matrix, so the actual density would be  somewhat lower, but the same method was used for all samples.  Layer (IV) is commonly known as the depleted layer and has been reported for ZrB2-SiC and HfB2-SiC systems.8,13-17,21 The average thickness of the depleted layer in the HS samples was 6 mm; however, samples HSWC and HSWB did not have a layer  (IV); the entire porous interface between the bulk and the oxide is comprised of HfO2. As in the 16001C oxidized samples, SiO2 scale in each of the samples had Al impurities whose concentration varied throughout the scale, but was never 41 mol%  the  assuming Al as the only impurity in SiO2. W-containing grains are present throughout all the layers in  both HSWC and HSWB. Examples of the W-containing grains are indicated by arrows in Fig. 5(d). EDS of the larger (1-2 mm)  grains  showed the presence of W and O. A third phase com prised of Hf, Si, and O is recognizable between HfO2 grains and the amorphous SiO2. The phase is likely HfSiO2 (labeled in Fig. 5(d)) as evidenced by its presence at 16001C and the increase  in the peak intensity observed in the XRD patterns of all the samples after oxidation to 18001C. Besides HfSiO4, only monoclinic and tetragonal HfO2 are present in the XRD patterns of all samples heated to 18001C.  (C)  Oxidation at 20001C:  Figure 6 is a series of com bined micrographs taken of the oxide scale resulting from exposure to 20001C. The layers of oxide scale for the HS sample heated to 20001C are: (I) a dense SiO2 outer layer with distributed HfO2 grains; (II) porous HfO2 penetrated by SiO2 (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C; and (IV) porous HfB2 with inclusions of different concentrations of Si, O, and C. The total scale thickness (Table I)  Table II.  Peak Shifts Observed for HfB2 and HfC in Samples HSWC and HSWB  Peak  Peak shift in 2y  HSWC  HSWB  HfB2 (001) HfB2 (100) HfB2 (101) HfB2 (002) HfB2 (110) HfC (111)  0.01  0.04  0.01  0.03  0.03  0.06  0.05  0.09  0.04  0.05  0.5  0.3  HfC (200)  0.5  0.3  HfC (220)  0.7  0.5  Table III.  Oxide Scale Thickness  Temperature (1C)  Oxide scale thickness (mm)  HS  HSWC  HSWB  1600  3573.2 7875.5 826756.1  3775.3 6775.5 537728.1  3472.3 7677.2 565723.8  1800  2000  Fig. 4. SEM micrograph of the W-containing phases that are found in the oxide scale after samples HSWC and HSWB are oxidized to 16001C.  Black arrows indicate the intergranular phase found between HfO2 that is composed of W-O and possibly Hf, while the white arrow indicates  spherical grains that are found throughout the oxide scale that also con tain W-O and possibly Hf.  August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2603  \\x0c', '2604  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  Fig. 5. SEM micrographs of the (a) HS, (b) HSWC, and (c) HSWB samples after oxidation at 18001C. The white arrows indicate the location of W containing grains in HSWC (d) while the HfSiO4, HfO2, and SiO2 phases are labeled. HfSiO4 was found in all three samples oxidized to 18001C. The oxide scale layers are indicated in (a) as (I) dense SiO2 outer layer with distributed HfO2 grains, (II) porous HfO2 penetrated by SiO2, (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C, and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C.  was  reduced by the addition of WC and WB while the SiO2 penetration into the scale and the density increased. In the  HSWC and HSWB samples  the distinction between layer  (I)  and layer  (II)  is blurred through the  extensive  formation of  HfSiO4. Layer (III) is reduced from an average thickness of 357 mm in HS to 170 mm in HSWC and 93 mm in HSWB. In the  Fig. 6. SEM micrographs of the (a) HSWB, (b) HSWC, and (c) HS samples after oxidation at 20001C. The layers: (I) a dense SiO2 outer layer with distributed HfO2 grains; (II) porous HfO2 penetrated by SiO2 and HfSiO4; (III) porous HfO2 with inclusions of varying concentrations of Si, O, and C; and (IV) porous HfB2 with inclusions of varying concentrations of Si, O, and C are indicated in the images. The magnified image in (a) shows the SiO2 and HfSiO4/HfO2 in layer (I) and (II), while the magnified image in (b) shows W-rich phases found in layer (III) as indicated by the arrows.  \\x0c', 'August 2011  Oxidation Resistance of Hafnium Diboride Ceramics with SiC, WB, or WC  2605  Fig. 7.  (a) SEM micrograph showing three distinct glass phases of the HSWB sample oxidized to 20001C: (A) W-rich, (B) SiO2 with Al, Fe, Ca, and Ba, (C) SiO2-rich, and (D) HfO2. (b) A TEM section from layer (I). The dark phase is W-rich. (c) TEM micrograph of a section that was prepared after the surface glass had been removed. The inset is the SEM micrograph of the surface of the sample showing the Pt cap applied for FIB cutting. The HfO2 (the light phase in the SEM and the dark phase in TEM) was monoclinic as determined by selected area electron diffraction (not shown). (d) TEM-EDS of  the glass phase showing the composition of the light and dark glasses in (c).  HSWC (Fig. 6(b)) and HSWB (Fig. 6(a)) the average thickness of the depleted later (layer (IV)) is 22 mm in HSWC and 32 mm in HSWB compared with 90 mm in HS. Image analysis calculated a  layer (III) density of 87%, 96%, and 97% in HS, HSWC, and  HSWB, respectively.  The SiO2 in layer (I) of the HS sample had impurities of Al, Ca, Ba, and Fe whose concentration varied spatially. Addition ally, the SiO2 in the oxide scale contains the same impurities (Al, Ca, Fe, Ba) as were found in the HS sample but were also shown  to form phases with W. Figure 6(a) is a magnified view of layers  (I) and (II). Extensive HfSiO4 formation is observed surrounding the lighter HfO2 grains. The HfSiO4 does not contain W, but small pockets (indicated by an arrow) of Hf-Si-W-O can be  found in layer  (II) of both the HSWC and HSWB samples.  Figure 6(b) is a magnified image of layer (III) showing the W containing spheres that were found scattered throughout in both  the HSWC and HSWB samples.  SEM and TEM images of  the W-containing phases  in the  glass are shown in Fig. 7. In Fig. 7(a) SEM-EDS suggests three  distinct glass phases:  (2) SiO2 with Al, Fe, Ca, and Ba (labeled B), and (3) SiO2-rich (labeled C). A TEM section from layer (I) shows a similar microstructure for  (1) W-rich (labeled A),  taken of the glass phase. Fitting the EDS pattern under the as sumption that all the impurities are oxides, the light glass contains o0.6 mol% of any impurity with HfO2 having the highest concentration. The dark glass contained Si, O, Hf, W, Ca, Al, Ba, Fe, and possibly Mg. The low intensity of the MgKa peak  and its overlap with an Hf M peak make Mg difficult to discern  in this  spectra. Selected area electron diffraction analysis  re vealed all phases in Fig. 7(b) and those in 7(c), exclusive of the  HfO2, to be amorphous. The XRD patterns of the sample surfaces after oxidation at 20001C reveal prominent HfSiO4 peaks in the HSWB sample, which correlates well with the observed microstructure. Approximately 60% of the top surface of the 20001C oxidized HSWB is  comprised of  the dense HfSiO4/Hf-Si-W-O phase while localized regions of dense HfSiO4/HfO2 comprise approximately 30% of the top surface observed in 20001C oxidized HSWC.  Although HfO2 would be expected to transform to tetragonal at 20001C, the slow cooling rate (51C/min) would allow conversion  to the monoclinic phase at lower temperatures and account for the small tetragonal peak (o15% of the height of the monoclinic (\\x00111) peak) observed for all the samples.  the multiphase glass with the black phase being W-rich. Figure  7(c) is an overview of a section that was prepared after the sur face glass had been removed. The inset is the SEM micrograph  IV.  Discussion  of the surface of the sample showing the Pt cap applied for FIB  After  sintering, no pure WB or WC phases were observed  cutting. The HfO2 phase in TEM) was monoclinic as determined by selected area  (the light phase in the SEM and the dark  electron diffraction (not  shown). TEM-EDS (Fig. 7(d)) were  in the HSWC and HSWB samples. The W was  incorporated  in the HfB2 and HfC matrix as evidenced by their peak shifts in the XRD spectra. HfC can be formed by a reaction of HfO2  \\x0c', 'and excess C. HfO2 is present as oxygen contamination on the HfB2 powders and oxidation during milling. Amorphous borides  (B2O3) also exist at from wear of the high-density  the surface. C can be introduced  polyethylene  bottles  during  milling or  from the graphite dies used to sinter  the powders.  The oxides may be consumed at elevated temperatures through  a reaction with free carbon by reactions (3) or (4) or with WC  (reaction [5]).  HfO2 ðsÞ þ B2O3 ðl Þ þ 5CðsÞ ! HfB2 ðsÞ þ 5COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x008:6 KJ=mol  (3)  HfO2 ðsÞ þ 3CðsÞ ! HfCðsÞ þ 2COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x00111:5 KJ=mol  (4)  HfO2 ðsÞ þ 3WCðsÞ ! HfCðsÞ þ 3WðsÞ þ 2COðgÞ DGr ð2026:85\\x0eCÞ ¼ \\x00489:3 KJ=mol  (5)  The presence of HfC after sintering diborides has been shown to be dependent on sintering temperature and additives.29,30  The phase diagram of HfB2 and WBB2 suggests a maximum 23 mol% at B23651C and W concentration of 10 mol% or above at temperatures above 12001C,31 while melting exper iments suggest  that approximately 4 mol% W can be incorpolattice.32 The same crystalline lattice and  rated into the HfB2 similar size allow 40 mol% of WC to be dissolved in HfC at 20271C.33 Limited literature exists for the HfC-WB system, but  the similar XRD peak shift values of HSWC and HSWB suggest  solubility of W in HfC in the HSWB sample. The W that  is  not incorporated into either HfC or HfB2 can be found in grains shown to contain W, O, C, and B. The addition of W-containing  phases  increases  the  sinterability  of  the  samples,  results  in smaller HfB2 SiC grains. A similar SiC microstructure was shown by Zhang et al.34 for pressureless sintered ZrB2-SiC with B4C additions and WC impurities. However, Chamberlain et al.24 do not show  grains,  and promotes  evolution of  acicular  the same SiC microstructure for WC impurities in hot pressed  ZrB2-SiC. It is possible that excess B or C (from WB, WC, or B4C34) in the system can promote the b (cubic) to a (hexagonal) transformation at temperatures below 18001C, with the a-SiC having acicular grain morphology.35 Unfortunately the SiC XRD peak has a very low intensity and the main b and a-SiC peaks are within 0.071, so phase identification is imprac tical. Engineering the microstructure may prove beneficial as it  has been shown that the size and morphology of HfB2 grains can affect strength.29,35,36 The oxidation of HS, HSWC, and HSWB samples resulted  the SiC and  in oxide scales with comparable total thicknesses and scale morphologies at 16001 and 18001C. The difference in the aver age thickness of the total oxide scale for each sample was within  the spread measured for an individual sample. The oxide scales formed on the HSWC and HSWB samples at 20001C were sig nificantly different  than the scale formed on HS. The HSWC  and HSWB samples had less porous scales that were 30% thin ner than the scale formed on the HS sample. The three methods  by which W-additions may impact oxidation resistance are (1)  controlling the phase transformations of the HfO2 scale, (2) increasing the viscosity of the SiO2 scale, or (3) promoting densification of the HfO2 scale. The addition of W neither promoted nor delayed the monoclinic to tetragonal phase transformation,  thus not affecting the density of the HfO2 by this method. Because no W was found in the SiO2 at 16001 and 18001C by SEM-EDS the addition of W-containing phases would not be  expected to alter  the viscosity or melting temperature of  the  SiO2 glass and the SiO2-rich layer would be expected to be protective at 16001 and 18001C.3,4,18,21 A two-phase glass is not 18001C which (17231C),  observed  until  above  is  above  the melting  temperature of pure SiO2 form as the glass cools. Borate and silicate glasses with Group  so the  two phase may  IV-VI transition metal (Hf, Zr, Ti, Ta, W) oxides tend to exhibit microphase separation and crystallization upon cooling.15,37,38  Glass  compositions  that  exhibit phase  separation posses  an  increased viscosity in the single-phase liquid.38  two-phase  region compared with a  Only a few phase diagrams for W and Hf in B2O3 or SiO2 are reported. The most studied diagram, HfO2-SiO2, reveals that at 20001C a single liquid phase exists for compositions with  approximately 60 mol% (or greater) of SiO2 and that at lower SiO2 concentrations the Si-Hf-O liquid phase is found in equilibrium with HfO2.27,37 In the WO3-B2O3 system a single-phase liquid is formed at all compositions above 14301C.39 No phase  diagram exists for HfO2-B2O3 but a study of glass melts found that o1 mol% HfO2 could be dissolved in B2O3.40 However, the same study found that sodium borosilicate or aluminum  borosilicate glasses can dissolve up to 17 mol% HfO2 with the allowable HfO2 concentration dependent on the concentration of the other glass species.39 As evidenced by Fig. 7 the glass chemistry found in the 20001C sample has spatial variances in  chemistry. The presence of Al, Na, Ca, Ba, and Fe in the glass  could originate from the impurities in the starting powders or  be incorporated from the Ca-stabilized ZrO2 crucible that holds the sample. Group I and Group II impurities have been  observed in our previous experiments and also reported in the literature,16 but have not received much attention to date. How ever,  as  the development of new UHTC compositions  seek  to add elements  in order  to improve glass properties  such as  viscosity the presence of these impurities may play a crucial role  in the outcome of these efforts.  The  phase  diagrams  for WO3 with SiO2 temperatures of interest, but  or  B2O3 inspection of  are  unavailable at  the  other  transition metal oxides with SiO2 such as Nb2O5-SiO2 and Cr2O3-SiO2 suggest two-phase glasses above 16951 and 22001C, respectively.41,42 Therefore the observed existence of a  two-phase  region  in  the WO3-SiO2 glass may improve oxidation resistance by the glass at 20001C and/or during  system seems  logical.  The  two-phase  increasing the viscosity of  its slow cool down.  The  impact of densification in the HfO2-based layers was examined in the HSWC and HSWB samples. The HfO2 contained W both as individual grains and as a phase between  HfO2 grains. Reactions (6)-(8) describe the possible oxidation of WC and WB.43,44  WCðcrÞ þ 5=2 O2 ðgÞ ! WO3 ðcrÞ þ CO2 ðgÞ  (6)  2 WBðcrÞ þ 9=2 O2 ðgÞ ! 2 WO3 ðcrÞ þ B2O3 ðl Þ  (7)  WO3 ðcrÞ ! WO3 ðgÞ  (8)  The HfO2-WO3 phase diagram shows a solubility of WO3 of about 5 mol% at 16001-20001C, while above 5 mol% the HfO2 solid solution will be accompanied by a liquid.22 The solubility  of W in HfB2 and HfC is between 4 and 40 mol% suggesting localized regions of W concentrations necessary for phase sep aration could exist. The resulting liquid has a melting point of 14731C. The presence of  this  liquid can promote liquid-phase  sintering of  the HfO2 scales for the HSWC and HSWB samples. The reduction in the length of the total SiC-depleted layer at 18001C in the samples  resulting in the observed denser oxide  containing W can be explained by a reduction in oxygen pen etration through the denser oxide  scale. Upon cooling,  the  resulting crystalline phases are a HfO2 solid solution with o5 mol% WO3 and WO3 with no solid solubility for HfO2. SEM-EDS confirmed the existence of W-rich phases between  the HfO2 grains in the postoxidation analysis. In addition to the denser HfO2 layer (layer (III)), layer (II) in the samples with W additions is composed of HfSiO4 and HfO2. Although a dense HfSiO4 would be expected to have a lower oxygen diffusion coefficient than the porous SiO2-filled HfO2 regions of HS, the decomposition  found  in  the  outermost  2606  Journal of the American Ceramic Society—Carney et al.  Vol. 94, No. 8  \\x0c', 'temperature of HfSiO4 is 17261C27 so that the HfSiO4 is likely formed on cooling and would not be responsible for the decrease in oxygen penetration at 20001C.  The  combined effect of  the more viscous outer  layer and  denser inner layer promote oxidation resistance and provide for an overall thinner scale at exposure temperatures of 20001C in  HSWC and HSWB.  V.  Conclusion  Dense HfB2-SiC samples were prepared with additives of WC and WB. Both the WC and WB additives promoted sintering  while reducing the grain size of HfB2. Solid solutions of W-Hf- B and W-Hf-C were formed in both samples. The WCand WB-containing HfB2-SiC samples oxidized to 16001 and 18001C exhibited similar oxide scales as the HfB2-SiC sample. However, the samples with WC and WB showed a 30% reduction in scale thickness when the samples were oxidized at 20001C  due to a more viscous phase separated glass found in the out ermost  regions  of  the  scale  and  a  denser  inner HfO2  that  restricted oxygen penetration to the sample.  Acknowledgment  The authors would like to thank Pavel Mogilevsky for his help in preparing the  TEM-FIB sample and useful discussions regarding TEM-EDS.  References  1R. L. Cloughtery, R. L. Pober, and L. Kaufman, ‘‘Thermogravimetric Study of the Oxidation of ZrB2 in the Temperature Range of 800-15001C,’’ Trans. Met. 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},{
  "_id": 191,
  "PDF": "Oxidation resistance of hafnium diboride-silicon carbide from 1400 to 2000  degrees C.pdf",
  "Text": "['Oxidation resistance of hafnium diboride—silicon carbide from 1400 to 2000 °C  Carmen M. Carney  Received: 1 April 2009 / Accepted: 5 August 2009 / Published online: 18 August 2009 Ó Springer Science+Business Media, LLC 2009  Abstract  Oxidation resistance tests were carried out on  HfB2-20 vol.% SiC prepared by spark plasma sintering. The dense samples were exposed from 1400 to 2000 °C in  an ambient atmosphere for 1 h. For comparison,  the same  material was tested using an arc jet  to simulate an atmo spheric reentry environment. The oxidation properties of  the samples were determined by measuring the weight gain  per unit surface area and the thicknesses of the oxide scale.  The oxide scale consists of a SiO2 outer layer, porous HfO2  layers, and an HfB2 layer depleted in SiC. A transition in  HfO2 morphology  from equixed  to  columnar and a 1900 °C  decrease  in  SiO2  viscosity  between  1800  and  accompanied  a  rapid  increase  in weight  gain  and  scale  thickness.  Introduction  Due  to their high melting temperatures,  transition metal  diborides such as ZrB2 and HfB2 commonly referred to as  ultra high-temperature ceramics  (UHTCs) are being con sidered as  candidates  for  leading edges of  sharp-bodied  reentry  vehicles  [1-3]. While ZrB2-based UHTCs  have  received a majority of the attention because of their lower 10.5 g/cm3  density  (6.09  vs.  for HfB2)  and  lower  cost,  HfB2-based materials have been shown to be more oxida tion  resistant  [4-6]. When HfB2  is  exposed  to  air  at  elevated temperatures, HfO2  and B2O3  are  formed. The  HfO2 scale is porous and the B2O3 evaporation limits the use of temperatures because of its 450 °C melting point [7]  and high vaporization pressure. For protection at higher  temperatures, HfB2  is  commonly mixed with SiC. The  addition of SiC allows the formation of SiO2 at elevated ([1100 °C).  temperatures  The  SiO2  reacts with B2O3  forming a borosilicate glass. After a series of studies in the  late 1960s,  the most commonly tested samples are HfB2  with 10-30 vol.% SiC [2, 8]. Most evaluations of UHTC  samples reported in the literature are limited to a maximum of 1630 °C by the air furnaces commonly used [2, 8-10]. It  is not known if the SiO2 offers sufficient protection at temperatures above 1630 °C, given that its melting point is 1713 °C [7]. For these materials to be used as leading edge  materials  on  hypersonic  aircraft  they must 2000 °C  be  able  to  withstand  temperatures  up  to  in  a  reentry  environment.  Oxidation at  the  lower  temperatures  (  \\x0c', 'necessary to utilize a test method such as high-temperature  furnace oxidation to properly evaluate materials prior  to  expensive arc jet  testing, and to study long-term oxidation  behavior.  In this  study,  the oxidation behavior of HfB2 20 vol.% SiC (hereafter HfB2-SiC) is studied during isotemperatures of 1400 to 2000 °C, (a)  thermal exposures at  to evaluate their behavior at temperatures greater than 1630 °C, (b) to compare them with data/behavior reported temperatures below 1630 °C, and (c) to  in the literature at  compare them with behavior during arc jet  testing. Weight  gain per surface area data are used as the measure of oxi dation resistance. Scanning electron microscopy and X-ray  diffraction  analysis  of  the  oxide  scale  are  used  to  gain  insight on the characteristics of the scale after exposure to  high-temperature regimes.  Experimental procedure  HfB2  (Cerac, Milwaukee Wisconsin)  and  SiC (Reade  Advanced Materials, East Providence Rhode Island) were  used to mix 80-vol.% HfB2 and 20-vol.% SiC. The b-SiC  was 45-55 nm powder  (reported) with 97.5% purity. The  HfB2 had a mean particle size of 4.6 lm (measured) and  99.5% purity.  Powders were  ball milled  using  Si3N4  grinding media in isopropanol for 18 h. The powders were  then dried at  room temperature while stirring followed by  18 h of dry milling with the same Si3N4. The weight loss of  the Si3N4 grinding media was 0.1 wt% based on the total  powder  weight.  The  powder  was  sieved  through  an  80-mesh screen before sintering.  A total of 170 g of the dried powder were loaded into a  40-mm graphite die coated with BN and lined with graphite  foil. The sample was sintered using spark plasma sintering  (SPS: FCT Systeme GmbH Model HPD 25-1, Rauenstein, Germany) with a heating and cooling rate of 50 °C/min and a maximum temperature of 2100 °C (achieved using 6 V  and 3.4 kA). The hold time was 17 mins. The temperature  was measured  by  an  optical  pyrometer focused on the punch *5 mm from the  bottom of  a  bore  hole  in  the  powder. A vacuum of 150 Pa was maintained for the entire  heating  cycle. A DC current with  a  pulse  sequence  of  10 ms on and 5 ms off with a single pulse was used for  heating. A uniaxial load of 32 MPa was applied on the heat up to 1600 °C and held during the remaining heating cycle.  A 25.4-mm diameter flat-faced model was cut from the  sintered puck for arc jet  testing using rotational electrical  discharge machining (EDM). The geometry of the samples  was dictated by the advanced heating facility (AHF) arc jet  at NASA Ames [11]. The face of the sample was polished  to 1 lm to remove oxidation caused by EDM machining  and provide consistent  starting surface finishes. From the  remaining puck, 4-mm 9 2.5-mm 9 2-mm samples were  cut using a diamond saw and polished on all  sides  to a  1-lm finish for oxidation tests.  HfB2-SiC specimens were exposed to air at tempera100 °C between 1600 °C for  tures  every  1400  and  1 h  using a horizontal MoSi2 resistance heated tube furnace. Samples were also heated every 100 °C between 1600 and 2000 °C using a ZrO2 element furnace (ZrF-25: Shinagawa  Refractories Co., Tokyo Japan). The heating and cooling rates used for both furnaces were 5 °C/min. The heating  rate was  limited by the zirconia heating element. All  the  samples were placed on a zirconia crucible so that  the two  sides with  the  largest  surface  area were  parallel  to  the  zirconia  crucible. The  zirconia  crucibles were  concave  such that  the samples only made two line contacts along  their corners with the crucible. Both the sample and the  zirconia crucible weights were measured before and after  oxidation using a balance with 0.01-mg precision. At  least  two samples were tested at each temperature.  The  flat-faced  models  were  exposed  to  sustained  enthalpy flows using the AHF arc jet at NASA Ames. The  main air mass flow rate during the test was 80 g/s and the  stagnation  pressure was  0.1 atm.  The  specimens were  located at  a distance of 6 cm from the  exit nozzle. The  cold-wall heat flux as measured using a copper slug 102260 W/cm2.  mm diameter  hemisphere  calorimeter was  This resulted in a surface temperature between 1450 and 1570 °C for 3.6 min as measured by the three pyrometers.  A  4-mm 9 2.5-mm 9 2-mm  rectangular  sample  was  heated in the furnace by pushing the sample into and out of  the heated furnace within approximately 1 min and holding the sample at 1500 °C for 4 min to provide a comparison  with the arc jet sample.  The microstructures  of  the  as  sintered  and  oxidized  samples were studied by polishing the samples to a 1-lm  finish. The furnace-heated samples were polished perpen dicular to the top and bottom faces, where the top face was  defined as exposed surface and the bottom face was  the  surface  facing the  zirconia  crucible. The  arc  jet  sample  was cut with a diamond-loaded saw blade approximately  9 mm from the  center  of  the  sample  and  polished  per pendicular  to  the  exposed  surface. The microstructures  were characterized using SEM (Quanta, FEI, Hillsborough  OR)  along with  energy  dispersive  spectroscopy  (EDS:  Pegasus 4000, EDAX, Mahwah NJ) for chemical analysis.  The crystallography analysis of  the oxide was performed  on an X-ray diffractometer with Cu Ka  radiation (XRD:  Rigaku 2500, Tokyo Japan). The density of  the samples  was measured  using  the Archimedes method  and He  pycnometry  (Accupyc  1330, Micromeritics,  Norcross,  GA). The reported density values are an average of seven  measurements.  5674  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'J Mater Sci  (2009) 44:5673-5681  5675  Results  Microstructure  He pycnometry analysis of the unprocessed HfB2-SiC powder showed a density of 9.28 g cm3. The sintered HfB2- SiC sample was found to have a density of 9.23 g cm3 by the Archimedes’ test method and 9.24 g cm3 by He pycnome try. The agreement between the He pycnometry and Archi medes’ values  suggest no open porosity in the sample as  expected and the individual values both give a density over  99%. Figure 1 shows a SEM micrograph of  the polished  surface. Although some SiC grain pullout was observed to  occur during polishing,  the microstructure is regular with  negligible residual porosity. The SiC, dark features, in Fig. 1  is  homogeneously  dispersed  inter-granularly within  the  HfB2 matrix. The individual SiC grains within the HfB2 matrix are on the order of 2-3 lm in diameter.  Weight gain due to oxidation (furnace heating)  The average weight gain per surface area as a function of  temperature is plotted in Fig. 2a. The total weight gain was  calculated by weighing both the sample and the zirconia  crucible before and after  the heating cycle and adding the  two weight-change  values. The  samples  heated  in  the  MoSi2 element  tube furnace are plotted with the samples  heated in the ZrO2 element furnace. Samples were heated at 1600 °C in both furnaces for comparison. The values are within 0.2 mg/cm2. Two samples were heated at each temperature. Only a small weight gain of 1.5 and 2.5 mg/cm2 is  observed at 1400 and 1500 C, respectively. The weight surface area at 1600 and 1700 °C increases  gain per  less  Fig. 2  a Weight gain per surface area as a function of temperature for  the MoSi2 the ZrO2 furnace-heated samples (open square). b The fraction of total weight  furnace-heated  samples  diamond)  (filled  and  gain that was measured for  the sample (-) and the crucible (open  square), the weight gain measured for the crucible resulted from SiO2 flow onto the substrate. c The thickness of the oxide scale formed on  the HfB2-SiC after exposure was measured in at least 12 places along the sample and the average (filled  the top (as defined in the text) of  diamond), minimum (-), and maximum (?) values are plotted with  respect  to temperature  than 0.5 mg/cm2 over the 1500 °C value. The weight gain 1900 °C more 1700 °C value.  doubles over  than  the  at  Fig. 1  SEM micrograph  showing  the microstructure  of  the  as  sintered HfB2-SiC sample. Some SiC grain pullout occurs during  polishing  Literature values for weight gain during oxidation in air for  HfB2-20 vol.% SiC without  sintering aids are rare. From  Opila et al. [6] the weight gain was calculated from the provided rate constant to be 1.6 mg/cm2 for samples held at 1627 °C for 1 h. This value is lower than the values found  in this study, but the use of slower heating rates effectively  increases  the  time  at  oxidizing  temperatures  compared  to the literature values. Licheri et al. [13] heated HfB2-26.5 vol.% SiC to 1450 °C at 2 °C/min and found  123  \\x0c', 'approximately 0.5-mg/cm2 weight gain upon heating and 0.5-mg/cm2 weight  gain  after  a  1-h  hold. Assuming  a  similar weight gain during cooling,  the values are compa rable with the data presented.  Post-oxidation observations suggested that some of  the  glass had flowed onto the zirconia crucible. To determine  the degree of glass flow,  the sample and the zirconia cru cible were weighed individually, except for one of the samples heated to 2000 °C where the sample adhered to  the crucible after heating. Figure 2b shows the fraction of  the overall weight gain attributable to the sample and to the  zirconia crucible. Since the zirconia crucible when heated  alone did not show any weight gain at  these temperatures,  the  change  in weight of  the crucible  is attributed to the  glass flow from the sample onto the crucible as observed after oxidation. Between 1400 and 1700 °C the sample was  responsible for 95% or greater of the total weight gain. The  weight change of the zirconia crucible was responsible for a majority of the weight change starting at 1800 °C (61%) and by 2000 °C accounted for 96% of  the weight gain.  The thickness of the total oxide scale is shown in Fig. 2c the temperatures except 2000 °C where the entire  for all  HfB2-SiC sample was oxidized. The total oxide scale is  taken to be any region of the sample that has been affected  by  oxygen  interaction.  The maximum and minimum  thicknesses measured are recorded along with the average  thickness to show the variation in the scale thickness. For  consistency, all  the measurements were made on the top  surface of the sample far enough away from the corners so  that  the more  rapid  oxidation  that  occurs  there  did  not  influence  the  results. At  least 12 separate measurements  were averaged for each sample. The scale thickness is less than 20-lm below 1500 °C and 1800 °C.  increases  to  65 lm at  Oxide scale morphology (furnace heating)  SEM micrographs of the samples heated between 1400 and 2000 °C in the static air  furnaces  for 1 h are depicted in  Fig. 3a-h.  In general,  four  layers  are defined (as  shown  in Fig. 3a-g) that will be used in the discussion. The layers  in each sample were identified by EDS analysis of the dark  (Si containing) and light (Hf containing) phases throughout  the entire oxide scale. Layer  I,  the topmost  layer, consists  of primarily Si and O (presumably amorphous SiO2). B2O3 has also been observed in this layer up to 1550 °C [5], but  this was not  tracked in this  study. Layer  II consists of a  lighter  phase  containing O and Hf,  and  a  darker phase  containing O and Si which was  interpreted as SiO2  infil trating porous HfO2. Layer  III  is predominantly Hf and O  (likely as HfO2) and Layer IV is predominately HfB2. Not  all  the four  layers are formed at each temperature. Their  thicknesses (Fig. 2c) and morphology (Fig. 3) were tracked  with temperature  to describe  the  evolution of  the oxide  scale with temperature. The following paragraphs describe  the evolution of  the layers.  Layer I (SiO2) does not completely cover the surface of the sample until 1600 °C. At 1600 °C the glassy layer between 10 and 20 lm in thickness. The layer grows with  is  increasing temperature 1900 °C (greater 2000 °C the SiO2  until  it  reaches  a maximum at  than  100 lm in  some  locations). At  thickness is less than 30 lm. Bubbles starting at 1500 °C and have  form in the glassy layer  a  dramatic impact on the morphology of the SiO2 layer at 1900 °C as seen in Fig. 3f. Additionally, grains shown to  contain Hf and O by EDS can be found in the SiO2 layer as 1400 °C and  low as  become more  frequent  at  higher  temperatures.  Layer II (porous HfO2 filled by SiO2) 1400 and 1500 °C. At  is nondescript at  these temperatures the SiO2 sits on  top of  the sample or only penetrates between voids at  the  surface, but reaches no further than one or two HfO2 grain thicknesses. The layer develops to about 10 lm at 1600 °C 1700 °C. At 1800 °C the  and  increases  in  thickness  at  morphology  of  the  layer  begins  to  change.  Instead  of  equiaxed HfO2 grains that  retain the original HfB2 micro structure,  the grains appear more columnar and the SiO2 scale at 1900 °C. At 2000 °C  fills a thicker porous HfO2  the HfO2 is distinctively columnar and Layer II has grown to around 300 lm.  Layer III, the predominantly HfO2 layer, exists between 1400 and 1700 °C. At 1400 and 1500 °C it comprises the the oxide scale (20-30 lm)  bulk of  lying just below the  SiO2 layer. At these temperatures the layer consists of HfO2  with inclusions of Si, C, and O as  shown by EDS. The  inclusions have varying concentrations of Si, C, and O and  thus will be referred to as Si-O-C. The inclusions reside in  the voids between the HfO2  that  are  the  same  size  and  morphology as the SiC grains within the bulk suggesting they are partially oxidized SiC. At 1600 and 1700 °C  that  Layer III is located below the layer of HfO2 filled by SiO2 and does not grow much more than 5 lm. Si-O-C incluII at 1600 and 1700 °C.  sions are also observed in Layer  Although  the  samples  presented  here  are mounted  in  a  C-containing epoxy, sister samples were polished using a  tripod polisher with no epoxy and the same Si-O-C inclu sions are found. The same Si-O-C inclusions were found by  the  author when  studying ZrB2-SiC samples  [14]  and  similar Si-O-C inclusions have been described in other ZrB2-SiC studies between 1400 and 1600 °C [15, 16]. At 1800 °C and above, Layer  III  is  replaced by Layer  IV,  the  predominately HfB2 layer. SiO2 fills the entire these temperatures. At 1800 °C the  porous HfO2  layer at  predominately HfB2 layer contains both SiO2 and Si-O-C layers at 1900 °C:  inclusions. Layer IV forms two distinct  HfB2 with SiO2 and HfB2 with Si-O-C inclusions (nearest  5676  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'J Mater Sci  (2009) 44:5673-5681  5677  Fig. 3  Series of SEM  micrographs showing the  progression of the oxide scale  from a to g 1400, 1500, 1600, 1700, 1800, 1900, and 2000 °C,  respectively. The dark area at  the top of each image is the  epoxy. The layers are defined  as:  (I) primarily SiO2, (II) SiO2  infiltrating HfO2, (III) HfO2  with inclusions of Si-O-C, and  (IV) HfB2 with Si-O-C and SiO2. At 1900 and 2000 °C, two  distinct  layers form within layer  (IV): HfB2 with SiO2 (IVa) and  HfB2 with Si-O-C (IVb). The  boundaries of the layers are  approximately outlined by the  white lines. The scale thickness  increases with temperature as  shown by the changes in scale  bars of the micrographs.  h Magnified region of the 2000 °C sample (g) showing  Layers II, IVb, and IVa  to the bulk). At 2000 °C Layer IV splits into three layers: a  layer of HfB2 with Si-O-C inclusions nearest  to the sur with Si-O-C inclusions which comprises the entire center 2000 °C,  the Si-O-C inclusions  are  of  the  sample. At  face, a layer of HfB2 with SiO2, and another layer of HfB2  almost entirely made up of C.  123  \\x0c', 'Oxide scale morphology (arc jet heating)  Figure 4a is an SEM micrograph of  the sample heated in  the arc jet with a hold time of 3.6 min. The hold temper ature seen on the face of the sample was between 1450 and 1570 °C. Figure 4b is 1500 °C in  a  furnace-heated sample  that was  heated  to  approximately  1 min  by  quickly  moving the sample into the heated region of  the furnace  and held for 4 min. The oxide scale thickness for the furthick (*7 lm com nace-heated sample is about half as to *18 lm)  pared  as  the  arc  jet  heated  sample. The  furnace-heated sample that was held at 1500 °C for 1 h had an average scale thickness of 20 lm. The morphology of  the layers for the arc jet and furnace-heated samples are the  same  and are mainly consists of porous HfO2. There  is  minimal penetration of  the SiO2  (bold arrows  in Fig. 4a)  into the HfO2, Layer II, and the predominately HfO2, Layer  III, with Si-C-O inclusions  (thin arrows  in Fig. 4a)  are  observed  in  both  the  samples.  Following  a  particular  grouping of SiC grains from the unaffected bulk to Layer II  (the primarily porous HfO2  layer) provides evidence that  the Si-O-C inclusions  are byproducts of SiC oxidation.  The Si-O-C fill  the same space once occupied by SiC and  the HfO2 maintains  the  structure  of  the  original HfB2  grains. These areas are boxed in Fig. 4b.  Oxide scale crystallography  In order to track the change in crystal structure of the oxide  scale, samples heated to 1600, 1700, 1800, 1900, and 2000 °C were analyzed using XRD. After cooling, XRD scans were taken for each of the samples between 15° and 80° 2 h. Figure 5 shows a region of interest that highlights in the sample heated to 1800 °C. The  the phases present  major peaks can be attributed to monoclinic HfO2 as expected. A small hump near 21° was observed in the XRD  scans  of  all  the  samples  before  background  subtraction  which  is  an  indicative  of  amorphous SiO2  [17]. Peaks  corresponding to tetragonal or cubic HfO2 between 1700 and 2000 °C. The peaks  are  found  for  the tetragonal  and cubic systems closely overlap and the determination  between the two structures is difficult especially when the  peaks are weak [18]. HfSiO4 is present in samples heated to 1600 and 1800 °C. HfSiO4 is the reaction product from the  combination of HfO2 and SiO2. Calculated HfO2-SiO2 phase diagrams show that HfSiO4 is stable up to *1726 °C  independent of the SiO2 mole fraction [19]. However, there  is limited thermodynamic data for  the HfO2-SiO2 system  and the experimental verification of the data mostly exists temperatures below 1000 °C [20-22].  for  Fig. 4  a SEM micrograph of the sample heated in the arc jet for 3.6 min at 260 MW/cm2 (1450-1570 °C). The oxide scale consists of  an HfO2 matrix with SiO2  in the top layer  (examples  indicated by  thick-lined  arrow)  and Si-O-C inclusions  (examples  indicated  by  thin-lined arrow). b SEM micrograph of the sample heated in the furnace for 4 min at 1500 °C. The regions where SiC crosses from the  unaffected bulk into Layer II are boxed. Note the magnification is 92  that  in (a)  Fig. 5 XRD after cooling of the sample oxidized in the ZrO2 furnace at 1800 °C showing the monoclinic HfO2  (^),  tetragonal HfO2  (*),  and HfSiO4 (?) phases  5678  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'Discussion  Oxidation  Figure 2  shows  the  relationship  between  oxide  scale  thickness, weight gain per surface area, and percent weight  gain of the sample versus the zirconia crucible. It is evident  from these plots  that  these properties are all  interrelated.  The weight gain per surface area (Fig. 2a) increases by less than 1 mg/100 °C from 1400 to 1800 °C until more than 4 mg from 1800 to 1900 °C. This  it  jumps  is  the same  temperature range that the SiO2-rich glass begins to flow more readily as expected from its 1713 °C melting point  and as evidenced by the increase in weight of the zirconia crucible (Fig. 2b). At 1900 °C,  the scale thickness  triples  over  the  values  found  for  the  lower  temperatures. This  increase in scale thickness correlates with the doubling of weight gain at 1900 °C. The turning point for the change in is at 1800 °C. Below this  the oxide scale structure  tem perature,  the oxide  scale  consists of well-defined layers  including an outer layer of SiO2-rich glass on top of HfO2  with some SiO2 penetration into the porous HfO2 grains  and in the more inner portions partially oxidized SiC in the form of Si-O-C inclusions. At 1800 °C,  the SiO2 pene trating HfO2 becomes the bulk of  the scale and a layer of  HfB2 with partially oxidized SiC is observed. HfSiO4 is observed after heating the sample to 1800 °C  (Fig. 5). The phase diagram of HfSiO4 predicts its stability to be *1726 °C and below. The HfSiO4 phase may be  formed  upon  heating  or  during  cooling.  Formation  of  HfSiO4 occurs when silicon atoms diffuse into the HfO2  crystal interstitially until the solubility limit is reached and has been shown to occur as low as 1100-1400 °C [20, 23].  This can occur  in both Layer  I between the HfO2 crystals  embedded in the SiO2 matrix and Layer  II when the SiO2  fills the HfO2 matrix. The formation of a crystalline solid  from an amorphous glass may help limit oxygen penetra tion and therefore limits oxide scale growth; however,  the  HfSiO4 phase is not continuous and there remains a sub stantial oxidation pathway through the  amorphous SiO2 rich glass. Additionally,  there does not seem to be a cor relation with the degree of oxidation and the presence of  HfSiO2.  The appearance of the tetragonal or cubic HfO2 peaks in the XRD was observed for samples heated to 1700 °C and  above making  it  likely  that  the  phase  transformation  is  monoclinic to tetragonal as expected from the Hf-O phase  diagram [19, 20]. This phase transformation from mono clinic to tetragonal HfO2 results in about 3.5% reduction in  volume [4]. The monoclinic to tetragonal transition may be  linked to the transition from equiaxed to columnar HfO2 grains as reported previously around 1650 °C [4, 24]. The  resulting volume decrease opens additional porosity in the  HfO2  matrix  aiding  the  increased  SiO2  penetration  observed  at  these  temperatures  and making  the  grains  appear more  elongated.  In  pure metal  diborides,  the  transformation from monoclinic to tetragonal is shown to greatly enhance the oxidation of HfB2 at *1700 °C [25,  26]. However,  the existence of SiO2 appears to delay this  effect since the total weight gain is less than 3% of the starting sample weight up to 1800 °C and doubles to 6% at 1900 °C.  A change in morphology may also be caused by a flow  of borosilicate glass  containing HfO2. Studies  involving  ZrB2-SiC [27, 28] discuss the convection of ZrO2-B2O3-  SiO2  liquid driven to the  surface by the volume change  associated with oxidation The ZrO2 then precipitates out of  the liquid as it reaches the vicinity of the surface when the  B2O3 evaporates. The metal oxide (MeO2) phase is pre dicted to precipitate due  to the  fact  that  its  solubility is  higher  in a borosilicate glass than in a pure silicate glass  [27]. The phase diagrams of HfO2-SiO2  [19] and ZrO2-  SiO2 [29] show that about 2 mol% ZrO2 or HfO2 may be  dissolved in the pure SiO2 liquid at the eutectic point (1687 °C for ZrO2-SiO2 and around 1450 °C for HfO2-  SiO2), while  the  amount of dissolvable MeO2 2000 °C approximately  increases  with  temperature. At  40 mol%  HfO2 or 11 mol% ZrO2 would be soluble in SiO2 before  precipitating  as  a  tetragonal  crystal. Upon  cooling,  the  excess HfO2 would precipitate in the SiO2 layer. The flow  of MeO2-B2O3-SiO2  liquid  is  predicted  to  leave  an  underlying  columnar MeO2  layer  and  smaller  equiaxed  oxide particles near or in the SiO2 layer. At 1400 and 1500 °C less than 25 distinct HfO2 grains are observed in  the SiO2 layer (Layer I) in the cross section of the samples.  These grains are identified by their contrast difference with SiO2 and confirmed with EDS. At 1600 °C the population of 1-2-lm HfO2 grains within the SiO2 2000 °C  layer  increases  (Fig. 3)  and  at  1900  and  the  grains  exist  throughout the SiO2 scale. A systematic oxidation study of ZrB2-SiC between 1600 and 1900 °C with hold times of  1 h [28] showed similar oxide scales as observed in this study. At 1600 °C, a SiO2 layer covers a ZrO2 layer whose  pores are filled with SiO2. A SiC-depleted region is not observed until 1700 °C, while oriented growth of ZrO2 is observed at 1800 °C. At 1800 °C,  the SiC-depleted region  is said to contain ‘‘minor’’ amounts of Si, but  there is no  discussion as to whether this is in the form of SiO2, SiC, or 1900 °C,  something  else. At  a  SiC-depleted  layer  is  observed in ZrB2  similar  to Layer IV found for 1800 °C and  the oxi dation  of HfB2-SiC at  higher.  Since  the  monoclinic to tetragonal transformation ZrO2 is between 1000 and 1150 °C,  temperature  of  the structural change  of ZrO2 from equiaxed to columnar between 1700 and 1800 °C would not be expected to be caused by a phase  transformation. Therefore,  the  columnar growth habit  in  J Mater Sci  (2009) 44:5673-5681  5679  123  \\x0c', 'both HfB2-SiC and ZrB2-SiC are from the interaction of  the MeO2 with the borosilicate glass and its convection. 2000 °C in the HfB2-SiC system  Between  1400 and  there are four competing factors contributing to oxidation:  (1) the formation of HfO2, (2) the formation and stability of  the SiO2  scale,  (3)  the presence of HfSiO4,  and (4)  the  transition in morphology and increase in porosity caused  by  a  phase  transformation  and/or  convection  of HfO2-  B2O3-SiO2  liquid.  From the  observations,  it  seems  as  though  the  stability  of  the  SiO2  layer  plays  the most  important  role  in oxidation protection with the  apparent  opening of porosity playing a supporting role. Methods that  have  been  proposed  to  increase  oxidation  resistance  include  engineering the glassy phase by either  inducing  phase separations by the addition of Group IV-VI  transi tion metals  [11] or by the addition of different combina tions  of  Na2O,  Al2O3,  and/or  B2O3  to  incorporate  additional HfO2 [9, 30]. However, decreasing the solubility  of HfO2 may aid in slowing the transport of HfO2 from the  porous  oxide  scale  to  the  SiO2  amorphous  scale  and  thereby  suppresses  the  formation  of  additional  porosity.  The incorporation of TaSi2 has been shown to increase oxidation resistance of ZrB2-SiC ceramics up to 1600 °C,  but has conflicting results at higher  temperatures  [6, 31].  Additionally, stabilizing agents may be alloyed within the  UHTC to delay or halt  the monoclinic to tetragonal phase  transformations of  the HfO2.  Comparison between testing variations  There is no significant difference between samples heated  in  the MoSi2  and ZrO2  furnaces. Comparisons  of  the  weight gain and morphology (not shown) of the samples heated to 1600 °C in both furnaces show the same results.  In  arc  jet  testing  one might  expect  a  catalytic  atom  recombination effect due to the presence of dissociated O  and N present  in the heating plasma of the arc jet (and also  at  the  leading  edges  in  hypersonic  flight);  however,  recombination effects at  the surface of ZrB2 and HfB2 and  particularly their oxides have been shown to be negligible  [32, 33]. Additionally,  the stagnation pressure inside the  arc  jet may  be  less  than  1 atm (in  this  test  stagnation  pressure was 0.1 atm) which can influence  the  active  to  passive transitions in the oxidation of  the samples. This is  especially  true with  the  oxidation  of SiC. When  active  oxidation of SiC occurs  the protective SiO2  layer can be  removed  as  gaseous  SiO.  This  active  oxidation may  account  for  the  very  thin  (0-0.8 lm)  outer  SiO2  layer  observed on the arc jet heated sample when compared to the *0.5-5-lm SiO2  layer  formed on the furnace-heated  sample. This  lack of outer SiO2 may lead to accelerated  oxidation and the longer oxide scale observed for the arc jet heated sample (*18 lm) compared to the sample that was  furnace heated for 4 min (*7 lm). Rezaie et al. [34] have shown that ZrB2-SiC heated to 1500 °C in a lower oxygen  partial pressure has a thinner outer SiO2 layer and a thicker  overall oxide  scale when compared to the  same  sample  heated in air. Some SiO2 may also be lost due to the shear  forces  of  the  plasma.  These  effects  would  become  increasing important with increases  in temperature as  the  SiO2 becomes more fluid. This effect does not change the morphology and chemistry of the oxide scale near 1500 °C  as the arc jet heated sample and the furnace-heated samples  held at 4 min and 1 h all consists of a SiO2 outer  layer  (Layer I) and a porous HfO2 layer with Si-O-C inclusions  (Layer II).  Arc jet tests published by Gasch et al. [11] with a steady-state temperature between 1690 and 1750 °C for a  total exposure time of 20 min can be used as a comparison for the high-temperature tests ([1600 °C) in this study. As in our 1600 and 1700 °C furnace-heated samples, the oxide  scale was described to consist mainly of a porous HfO2  layer. They also report an HfB2  layer with SiC removed  (there is no mention of Si-O-C inclusions in the descrip tion of the arc jet sample). The total oxide scale thickness  of the arc jet heated sample was 72 lm when compared to the 39-lm total oxide scale observed in the 1700 °C fur nace-heated sample of this study that was held at time. From our 1500 °C data  temper ature for a longer  and the  literature data, one can conclude that  the furnace tests can  reproduce the basic microstructures of the oxidized HfB2-  SiC found in arc jet  tests, but  the total  scale thickness  is  limited  by  a  thicker  protective SiO2  outer  layer  in  the  furnace tests.  Conclusion  Samples were heated in air  from 1400 to 2000 °C using a  MoSi2 and ZrO2  furnace. These  samples were  shown to  develop four layers over this temperature range. Below 1800 °C, the oxide scale consists of three layers: Layer I, a  protective  SiO2; Layer  II,  SiO2  infiltrating  into  porous  HfO2; and Layer III, predominantly porous HfO2. At 1800 °C and above, the oxide layer consists of three layers:  Layer  I, a protective SiO2; Layer  II, SiO2 infiltrating into  porous HfO2; and Layer IV, predominately HfB2. Regions  of the predominately HfO2 and HfB2 layers both contained  inclusions of Si-O-C indicating partially oxidized SiC. Up to 2000 °C,  the critical  factors influencing oxidation were  shown to be the flow of SiO2 glass and an apparent increase  in porosity of the HfO2 oxide scale. The outer oxide scale  must be protective against oxygen penetration and as shown  by the  comparison of  furnace-heated samples  to arc  jet  heated samples must be  resistant  to active oxidation.  In  order to increase the working times at elevated temperatures  5680  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', '([1800 °C)  it  is necessary to develop a more robust pro tective glassy scale through the use of additives with SiO2 to  increase the viscosity or melting temperature or to stabilize  the HfO2 crystalline phase.  Acknowledgements  This study was supported in part by the United  States Air Force Contract # FA8650-04-D-5233 with project manager  Michael Cinibulk. We acknowledge NASA-SCAP for  their critical  financial support of the arc jet operational capability at Ames. And we  would like to thank Sylvia Johnson and Matthew Gasch at NASA  Ames for  their assistance in performing the arc jet  testing.  References  1. Monteverde F, Scatteia L (2007) J Am Ceram Soc 90(4):130  2. Fenter JR (1971) SAMPE Quart 2(3):1  3. Fahrenholtz WG, Hilmas GE, Talmy  IG, Zaykoski A (2007)  J Am Ceram Soc 90(5):1347  4. Kaufman L, Clougherty EV, Berkowitz-Mattuck JB (1967) Trans  Metall Soc AIME 239:458  5. Hinze  JW, Tripp WC, Graham HC (1975)  J Electrochem Soc  122(9):1249  6. Opila E, Levine S, Lorincz J  (2004)  J Mater Sci 39:5969. doi:  10.1023/B:JMSC.0000041693.32531.d1  7. Lide DR (1996) CRC handbook of chemistry and physics, 77th  edn. CRC Press Inc, Boca Raton (Sect. 4)  8. Clougherty EV, Pober RL, Kaufman L (1968) Trans Metall Soc  AIME 242:1077  9. Lespade P, Richet N, Goursat P (2007) Acta Astronautica 60:858  10. Monteverde F (2005) Corr Sci 47:2020  11. 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Bongiorno A, Fo¨ rst CJ, Kalia RK, Li J, Marschall J, Nakano A  et al (2006) Mater Res Soc Bull 31:410  23. Nguyen QGN, Opila EJ, Robinson RC (2004) J Electrochem Soc  151(10):B558  24. Berkowitz-Mattuck JB (1966) J Electrochem Soc 113(9):908  25. Parthasarathy TA, Rapp RA, Opeka M, Kerans RJ (2009) J Am  Ceram Soc 92(5):1079  26. Talmy IG, Zaykoski JA, Opeka MM, Dallik S (2001)  In: Opila  EJ, McNallan MJ, Shores DA, Shifler DA (eds) High-temperature  corrosion  and materials  chemistry  III.  The  Electrochemical  Society, Pennington, p 144  27. Karldottir SN, Halloran JW (2008) J Am Ceram Soc 91(1):272  28. Zhang X-H, Hu P, Han J-C (2008) J Mater Res 23(7):1961  29. Butterman WC, Foster WR (1967) Am Min 52(5-6):880  30. Davis LL, Li L, Darab JG, Li H, Strachan D (1999) Mater Res  Soc Symp Proc 556:313  31. Opeka MM, Talmy IG, Zaykoski JA (2004) J Mater Sci 39:5887.  doi:10.1023/B:JMSC.0000041686.21788.77  32. Marschall  J, Chamberlain A, Crunkleton D, Rogers B (2004)  J Spacecr Rocket 41(4):576  33. Scatteia L, Borrelli R, Cosentino G, Beˆ che E, Sans  J-L, Balat Pichelin M (2006) J Spacecr Rocket 43(5):1004  34. Rezaie A, Fahrenholtz WG, Hilmas GE (2006) J Am Ceram Soc  89(10):3240  J Mater Sci  (2009) 44:5673-5681  5681  123  \\x0c']"
},{
  "_id": 192,
  "PDF": "Oxidation resistance of hafnium diboride-silicon carbide from 1400 to 2000 A degrees C.pdf",
  "Text": "['Oxidation resistance of hafnium diboride—silicon carbide from 1400 to 2000 °C  Carmen M. Carney  Received: 1 April 2009 / Accepted: 5 August 2009 / Published online: 18 August 2009 Ó Springer Science+Business Media, LLC 2009  Abstract  Oxidation resistance tests were carried out on  HfB2-20 vol.% SiC prepared by spark plasma sintering. The dense samples were exposed from 1400 to 2000 °C in  an ambient atmosphere for 1 h. For comparison,  the same  material was tested using an arc jet  to simulate an atmo spheric reentry environment. The oxidation properties of  the samples were determined by measuring the weight gain  per unit surface area and the thicknesses of the oxide scale.  The oxide scale consists of a SiO2 outer layer, porous HfO2  layers, and an HfB2 layer depleted in SiC. A transition in  HfO2 morphology  from equixed  to  columnar and a 1900 °C  decrease  in  SiO2  viscosity  between  1800  and  accompanied  a  rapid  increase  in weight  gain  and  scale  thickness.  Introduction  Due  to their high melting temperatures,  transition metal  diborides such as ZrB2 and HfB2 commonly referred to as  ultra high-temperature ceramics  (UHTCs) are being con sidered as  candidates  for  leading edges of  sharp-bodied  reentry  vehicles  [1-3]. While ZrB2-based UHTCs  have  received a majority of the attention because of their lower 10.5 g/cm3  density  (6.09  vs.  for HfB2)  and  lower  cost,  HfB2-based materials have been shown to be more oxida tion  resistant  [4-6]. When HfB2  is  exposed  to  air  at  elevated temperatures, HfO2  and B2O3  are  formed. The  HfO2 scale is porous and the B2O3 evaporation limits the use of temperatures because of its 450 °C melting point [7]  and high vaporization pressure. For protection at higher  temperatures, HfB2  is  commonly mixed with SiC. The  addition of SiC allows the formation of SiO2 at elevated ([1100 °C).  temperatures  The  SiO2  reacts with B2O3  forming a borosilicate glass. After a series of studies in the  late 1960s,  the most commonly tested samples are HfB2  with 10-30 vol.% SiC [2, 8]. Most evaluations of UHTC  samples reported in the literature are limited to a maximum of 1630 °C by the air furnaces commonly used [2, 8-10]. It  is not known if the SiO2 offers sufficient protection at temperatures above 1630 °C, given that its melting point is 1713 °C [7]. For these materials to be used as leading edge  materials  on  hypersonic  aircraft  they must 2000 °C  be  able  to  withstand  temperatures  up  to  in  a  reentry  environment.  Oxidation at  the  lower  temperatures  (  \\x0c', 'necessary to utilize a test method such as high-temperature  furnace oxidation to properly evaluate materials prior  to  expensive arc jet  testing, and to study long-term oxidation  behavior.  In this  study,  the oxidation behavior of HfB2 20 vol.% SiC (hereafter HfB2-SiC) is studied during isotemperatures of 1400 to 2000 °C, (a)  thermal exposures at  to evaluate their behavior at temperatures greater than 1630 °C, (b) to compare them with data/behavior reported temperatures below 1630 °C, and (c) to  in the literature at  compare them with behavior during arc jet  testing. Weight  gain per surface area data are used as the measure of oxi dation resistance. Scanning electron microscopy and X-ray  diffraction  analysis  of  the  oxide  scale  are  used  to  gain  insight on the characteristics of the scale after exposure to  high-temperature regimes.  Experimental procedure  HfB2  (Cerac, Milwaukee Wisconsin)  and  SiC (Reade  Advanced Materials, East Providence Rhode Island) were  used to mix 80-vol.% HfB2 and 20-vol.% SiC. The b-SiC  was 45-55 nm powder  (reported) with 97.5% purity. The  HfB2 had a mean particle size of 4.6 lm (measured) and  99.5% purity.  Powders were  ball milled  using  Si3N4  grinding media in isopropanol for 18 h. The powders were  then dried at  room temperature while stirring followed by  18 h of dry milling with the same Si3N4. The weight loss of  the Si3N4 grinding media was 0.1 wt% based on the total  powder  weight.  The  powder  was  sieved  through  an  80-mesh screen before sintering.  A total of 170 g of the dried powder were loaded into a  40-mm graphite die coated with BN and lined with graphite  foil. The sample was sintered using spark plasma sintering  (SPS: FCT Systeme GmbH Model HPD 25-1, Rauenstein, Germany) with a heating and cooling rate of 50 °C/min and a maximum temperature of 2100 °C (achieved using 6 V  and 3.4 kA). The hold time was 17 mins. The temperature  was measured  by  an  optical  pyrometer focused on the punch *5 mm from the  bottom of  a  bore  hole  in  the  powder. A vacuum of 150 Pa was maintained for the entire  heating  cycle. A DC current with  a  pulse  sequence  of  10 ms on and 5 ms off with a single pulse was used for  heating. A uniaxial load of 32 MPa was applied on the heat up to 1600 °C and held during the remaining heating cycle.  A 25.4-mm diameter flat-faced model was cut from the  sintered puck for arc jet  testing using rotational electrical  discharge machining (EDM). The geometry of the samples  was dictated by the advanced heating facility (AHF) arc jet  at NASA Ames [11]. The face of the sample was polished  to 1 lm to remove oxidation caused by EDM machining  and provide consistent  starting surface finishes. From the  remaining puck, 4-mm 9 2.5-mm 9 2-mm samples were  cut using a diamond saw and polished on all  sides  to a  1-lm finish for oxidation tests.  HfB2-SiC specimens were exposed to air at tempera100 °C between 1600 °C for  tures  every  1400  and  1 h  using a horizontal MoSi2 resistance heated tube furnace. Samples were also heated every 100 °C between 1600 and 2000 °C using a ZrO2 element furnace (ZrF-25: Shinagawa  Refractories Co., Tokyo Japan). The heating and cooling rates used for both furnaces were 5 °C/min. The heating  rate was  limited by the zirconia heating element. All  the  samples were placed on a zirconia crucible so that  the two  sides with  the  largest  surface  area were  parallel  to  the  zirconia  crucible. The  zirconia  crucibles were  concave  such that  the samples only made two line contacts along  their corners with the crucible. Both the sample and the  zirconia crucible weights were measured before and after  oxidation using a balance with 0.01-mg precision. At  least  two samples were tested at each temperature.  The  flat-faced  models  were  exposed  to  sustained  enthalpy flows using the AHF arc jet at NASA Ames. The  main air mass flow rate during the test was 80 g/s and the  stagnation  pressure was  0.1 atm.  The  specimens were  located at  a distance of 6 cm from the  exit nozzle. The  cold-wall heat flux as measured using a copper slug 102260 W/cm2.  mm diameter  hemisphere  calorimeter was  This resulted in a surface temperature between 1450 and 1570 °C for 3.6 min as measured by the three pyrometers.  A  4-mm 9 2.5-mm 9 2-mm  rectangular  sample  was  heated in the furnace by pushing the sample into and out of  the heated furnace within approximately 1 min and holding the sample at 1500 °C for 4 min to provide a comparison  with the arc jet sample.  The microstructures  of  the  as  sintered  and  oxidized  samples were studied by polishing the samples to a 1-lm  finish. The furnace-heated samples were polished perpen dicular to the top and bottom faces, where the top face was  defined as exposed surface and the bottom face was  the  surface  facing the  zirconia  crucible. The  arc  jet  sample  was cut with a diamond-loaded saw blade approximately  9 mm from the  center  of  the  sample  and  polished  per pendicular  to  the  exposed  surface. The microstructures  were characterized using SEM (Quanta, FEI, Hillsborough  OR)  along with  energy  dispersive  spectroscopy  (EDS:  Pegasus 4000, EDAX, Mahwah NJ) for chemical analysis.  The crystallography analysis of  the oxide was performed  on an X-ray diffractometer with Cu Ka  radiation (XRD:  Rigaku 2500, Tokyo Japan). The density of  the samples  was measured  using  the Archimedes method  and He  pycnometry  (Accupyc  1330, Micromeritics,  Norcross,  GA). The reported density values are an average of seven  measurements.  5674  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'J Mater Sci  (2009) 44:5673-5681  5675  Results  Microstructure  He pycnometry analysis of the unprocessed HfB2-SiC powder showed a density of 9.28 g cm3. The sintered HfB2- SiC sample was found to have a density of 9.23 g cm3 by the Archimedes’ test method and 9.24 g cm3 by He pycnome try. The agreement between the He pycnometry and Archi medes’ values  suggest no open porosity in the sample as  expected and the individual values both give a density over  99%. Figure 1 shows a SEM micrograph of  the polished  surface. Although some SiC grain pullout was observed to  occur during polishing,  the microstructure is regular with  negligible residual porosity. The SiC, dark features, in Fig. 1  is  homogeneously  dispersed  inter-granularly within  the  HfB2 matrix. The individual SiC grains within the HfB2 matrix are on the order of 2-3 lm in diameter.  Weight gain due to oxidation (furnace heating)  The average weight gain per surface area as a function of  temperature is plotted in Fig. 2a. The total weight gain was  calculated by weighing both the sample and the zirconia  crucible before and after  the heating cycle and adding the  two weight-change  values. The  samples  heated  in  the  MoSi2 element  tube furnace are plotted with the samples  heated in the ZrO2 element furnace. Samples were heated at 1600 °C in both furnaces for comparison. The values are within 0.2 mg/cm2. Two samples were heated at each temperature. Only a small weight gain of 1.5 and 2.5 mg/cm2 is  observed at 1400 and 1500 C, respectively. The weight surface area at 1600 and 1700 °C increases  gain per  less  Fig. 2  a Weight gain per surface area as a function of temperature for  the MoSi2 the ZrO2 furnace-heated samples (open square). b The fraction of total weight  furnace-heated  samples  diamond)  (filled  and  gain that was measured for  the sample (-) and the crucible (open  square), the weight gain measured for the crucible resulted from SiO2 flow onto the substrate. c The thickness of the oxide scale formed on  the HfB2-SiC after exposure was measured in at least 12 places along the sample and the average (filled  the top (as defined in the text) of  diamond), minimum (-), and maximum (?) values are plotted with  respect  to temperature  than 0.5 mg/cm2 over the 1500 °C value. The weight gain 1900 °C more 1700 °C value.  doubles over  than  the  at  Fig. 1  SEM micrograph  showing  the microstructure  of  the  as  sintered HfB2-SiC sample. Some SiC grain pullout occurs during  polishing  Literature values for weight gain during oxidation in air for  HfB2-20 vol.% SiC without  sintering aids are rare. From  Opila et al. [6] the weight gain was calculated from the provided rate constant to be 1.6 mg/cm2 for samples held at 1627 °C for 1 h. This value is lower than the values found  in this study, but the use of slower heating rates effectively  increases  the  time  at  oxidizing  temperatures  compared  to the literature values. Licheri et al. [13] heated HfB2-26.5 vol.% SiC to 1450 °C at 2 °C/min and found  123  \\x0c', 'approximately 0.5-mg/cm2 weight gain upon heating and 0.5-mg/cm2 weight  gain  after  a  1-h  hold. Assuming  a  similar weight gain during cooling,  the values are compa rable with the data presented.  Post-oxidation observations suggested that some of  the  glass had flowed onto the zirconia crucible. To determine  the degree of glass flow,  the sample and the zirconia cru cible were weighed individually, except for one of the samples heated to 2000 °C where the sample adhered to  the crucible after heating. Figure 2b shows the fraction of  the overall weight gain attributable to the sample and to the  zirconia crucible. Since the zirconia crucible when heated  alone did not show any weight gain at  these temperatures,  the  change  in weight of  the crucible  is attributed to the  glass flow from the sample onto the crucible as observed after oxidation. Between 1400 and 1700 °C the sample was  responsible for 95% or greater of the total weight gain. The  weight change of the zirconia crucible was responsible for a majority of the weight change starting at 1800 °C (61%) and by 2000 °C accounted for 96% of  the weight gain.  The thickness of the total oxide scale is shown in Fig. 2c the temperatures except 2000 °C where the entire  for all  HfB2-SiC sample was oxidized. The total oxide scale is  taken to be any region of the sample that has been affected  by  oxygen  interaction.  The maximum and minimum  thicknesses measured are recorded along with the average  thickness to show the variation in the scale thickness. For  consistency, all  the measurements were made on the top  surface of the sample far enough away from the corners so  that  the more  rapid  oxidation  that  occurs  there  did  not  influence  the  results. At  least 12 separate measurements  were averaged for each sample. The scale thickness is less than 20-lm below 1500 °C and 1800 °C.  increases  to  65 lm at  Oxide scale morphology (furnace heating)  SEM micrographs of the samples heated between 1400 and 2000 °C in the static air  furnaces  for 1 h are depicted in  Fig. 3a-h.  In general,  four  layers  are defined (as  shown  in Fig. 3a-g) that will be used in the discussion. The layers  in each sample were identified by EDS analysis of the dark  (Si containing) and light (Hf containing) phases throughout  the entire oxide scale. Layer  I,  the topmost  layer, consists  of primarily Si and O (presumably amorphous SiO2). B2O3 has also been observed in this layer up to 1550 °C [5], but  this was not  tracked in this  study. Layer  II consists of a  lighter  phase  containing O and Hf,  and  a  darker phase  containing O and Si which was  interpreted as SiO2  infil trating porous HfO2. Layer  III  is predominantly Hf and O  (likely as HfO2) and Layer IV is predominately HfB2. Not  all  the four  layers are formed at each temperature. Their  thicknesses (Fig. 2c) and morphology (Fig. 3) were tracked  with temperature  to describe  the  evolution of  the oxide  scale with temperature. The following paragraphs describe  the evolution of  the layers.  Layer I (SiO2) does not completely cover the surface of the sample until 1600 °C. At 1600 °C the glassy layer between 10 and 20 lm in thickness. The layer grows with  is  increasing temperature 1900 °C (greater 2000 °C the SiO2  until  it  reaches  a maximum at  than  100 lm in  some  locations). At  thickness is less than 30 lm. Bubbles starting at 1500 °C and have  form in the glassy layer  a  dramatic impact on the morphology of the SiO2 layer at 1900 °C as seen in Fig. 3f. Additionally, grains shown to  contain Hf and O by EDS can be found in the SiO2 layer as 1400 °C and  low as  become more  frequent  at  higher  temperatures.  Layer II (porous HfO2 filled by SiO2) 1400 and 1500 °C. At  is nondescript at  these temperatures the SiO2 sits on  top of  the sample or only penetrates between voids at  the  surface, but reaches no further than one or two HfO2 grain thicknesses. The layer develops to about 10 lm at 1600 °C 1700 °C. At 1800 °C the  and  increases  in  thickness  at  morphology  of  the  layer  begins  to  change.  Instead  of  equiaxed HfO2 grains that  retain the original HfB2 micro structure,  the grains appear more columnar and the SiO2 scale at 1900 °C. At 2000 °C  fills a thicker porous HfO2  the HfO2 is distinctively columnar and Layer II has grown to around 300 lm.  Layer III, the predominantly HfO2 layer, exists between 1400 and 1700 °C. At 1400 and 1500 °C it comprises the the oxide scale (20-30 lm)  bulk of  lying just below the  SiO2 layer. At these temperatures the layer consists of HfO2  with inclusions of Si, C, and O as  shown by EDS. The  inclusions have varying concentrations of Si, C, and O and  thus will be referred to as Si-O-C. The inclusions reside in  the voids between the HfO2  that  are  the  same  size  and  morphology as the SiC grains within the bulk suggesting they are partially oxidized SiC. At 1600 and 1700 °C  that  Layer III is located below the layer of HfO2 filled by SiO2 and does not grow much more than 5 lm. Si-O-C incluII at 1600 and 1700 °C.  sions are also observed in Layer  Although  the  samples  presented  here  are mounted  in  a  C-containing epoxy, sister samples were polished using a  tripod polisher with no epoxy and the same Si-O-C inclu sions are found. The same Si-O-C inclusions were found by  the  author when  studying ZrB2-SiC samples  [14]  and  similar Si-O-C inclusions have been described in other ZrB2-SiC studies between 1400 and 1600 °C [15, 16]. At 1800 °C and above, Layer  III  is  replaced by Layer  IV,  the  predominately HfB2 layer. SiO2 fills the entire these temperatures. At 1800 °C the  porous HfO2  layer at  predominately HfB2 layer contains both SiO2 and Si-O-C layers at 1900 °C:  inclusions. Layer IV forms two distinct  HfB2 with SiO2 and HfB2 with Si-O-C inclusions (nearest  5676  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'J Mater Sci  (2009) 44:5673-5681  5677  Fig. 3  Series of SEM  micrographs showing the  progression of the oxide scale  from a to g 1400, 1500, 1600, 1700, 1800, 1900, and 2000 °C,  respectively. The dark area at  the top of each image is the  epoxy. The layers are defined  as:  (I) primarily SiO2, (II) SiO2  infiltrating HfO2, (III) HfO2  with inclusions of Si-O-C, and  (IV) HfB2 with Si-O-C and SiO2. At 1900 and 2000 °C, two  distinct  layers form within layer  (IV): HfB2 with SiO2 (IVa) and  HfB2 with Si-O-C (IVb). The  boundaries of the layers are  approximately outlined by the  white lines. The scale thickness  increases with temperature as  shown by the changes in scale  bars of the micrographs.  h Magnified region of the 2000 °C sample (g) showing  Layers II, IVb, and IVa  to the bulk). At 2000 °C Layer IV splits into three layers: a  layer of HfB2 with Si-O-C inclusions nearest  to the sur with Si-O-C inclusions which comprises the entire center 2000 °C,  the Si-O-C inclusions  are  of  the  sample. At  face, a layer of HfB2 with SiO2, and another layer of HfB2  almost entirely made up of C.  123  \\x0c', 'Oxide scale morphology (arc jet heating)  Figure 4a is an SEM micrograph of  the sample heated in  the arc jet with a hold time of 3.6 min. The hold temper ature seen on the face of the sample was between 1450 and 1570 °C. Figure 4b is 1500 °C in  a  furnace-heated sample  that was  heated  to  approximately  1 min  by  quickly  moving the sample into the heated region of  the furnace  and held for 4 min. The oxide scale thickness for the furthick (*7 lm com nace-heated sample is about half as to *18 lm)  pared  as  the  arc  jet  heated  sample. The  furnace-heated sample that was held at 1500 °C for 1 h had an average scale thickness of 20 lm. The morphology of  the layers for the arc jet and furnace-heated samples are the  same  and are mainly consists of porous HfO2. There  is  minimal penetration of  the SiO2  (bold arrows  in Fig. 4a)  into the HfO2, Layer II, and the predominately HfO2, Layer  III, with Si-C-O inclusions  (thin arrows  in Fig. 4a)  are  observed  in  both  the  samples.  Following  a  particular  grouping of SiC grains from the unaffected bulk to Layer II  (the primarily porous HfO2  layer) provides evidence that  the Si-O-C inclusions  are byproducts of SiC oxidation.  The Si-O-C fill  the same space once occupied by SiC and  the HfO2 maintains  the  structure  of  the  original HfB2  grains. These areas are boxed in Fig. 4b.  Oxide scale crystallography  In order to track the change in crystal structure of the oxide  scale, samples heated to 1600, 1700, 1800, 1900, and 2000 °C were analyzed using XRD. After cooling, XRD scans were taken for each of the samples between 15° and 80° 2 h. Figure 5 shows a region of interest that highlights in the sample heated to 1800 °C. The  the phases present  major peaks can be attributed to monoclinic HfO2 as expected. A small hump near 21° was observed in the XRD  scans  of  all  the  samples  before  background  subtraction  which  is  an  indicative  of  amorphous SiO2  [17]. Peaks  corresponding to tetragonal or cubic HfO2 between 1700 and 2000 °C. The peaks  are  found  for  the tetragonal  and cubic systems closely overlap and the determination  between the two structures is difficult especially when the  peaks are weak [18]. HfSiO4 is present in samples heated to 1600 and 1800 °C. HfSiO4 is the reaction product from the  combination of HfO2 and SiO2. Calculated HfO2-SiO2 phase diagrams show that HfSiO4 is stable up to *1726 °C  independent of the SiO2 mole fraction [19]. However, there  is limited thermodynamic data for  the HfO2-SiO2 system  and the experimental verification of the data mostly exists temperatures below 1000 °C [20-22].  for  Fig. 4  a SEM micrograph of the sample heated in the arc jet for 3.6 min at 260 MW/cm2 (1450-1570 °C). The oxide scale consists of  an HfO2 matrix with SiO2  in the top layer  (examples  indicated by  thick-lined  arrow)  and Si-O-C inclusions  (examples  indicated  by  thin-lined arrow). b SEM micrograph of the sample heated in the furnace for 4 min at 1500 °C. The regions where SiC crosses from the  unaffected bulk into Layer II are boxed. Note the magnification is 92  that  in (a)  Fig. 5 XRD after cooling of the sample oxidized in the ZrO2 furnace at 1800 °C showing the monoclinic HfO2  (^),  tetragonal HfO2  (*),  and HfSiO4 (?) phases  5678  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', 'Discussion  Oxidation  Figure 2  shows  the  relationship  between  oxide  scale  thickness, weight gain per surface area, and percent weight  gain of the sample versus the zirconia crucible. It is evident  from these plots  that  these properties are all  interrelated.  The weight gain per surface area (Fig. 2a) increases by less than 1 mg/100 °C from 1400 to 1800 °C until more than 4 mg from 1800 to 1900 °C. This  it  jumps  is  the same  temperature range that the SiO2-rich glass begins to flow more readily as expected from its 1713 °C melting point  and as evidenced by the increase in weight of the zirconia crucible (Fig. 2b). At 1900 °C,  the scale thickness  triples  over  the  values  found  for  the  lower  temperatures. This  increase in scale thickness correlates with the doubling of weight gain at 1900 °C. The turning point for the change in is at 1800 °C. Below this  the oxide scale structure  tem perature,  the oxide  scale  consists of well-defined layers  including an outer layer of SiO2-rich glass on top of HfO2  with some SiO2 penetration into the porous HfO2 grains  and in the more inner portions partially oxidized SiC in the form of Si-O-C inclusions. At 1800 °C,  the SiO2 pene trating HfO2 becomes the bulk of  the scale and a layer of  HfB2 with partially oxidized SiC is observed. HfSiO4 is observed after heating the sample to 1800 °C  (Fig. 5). The phase diagram of HfSiO4 predicts its stability to be *1726 °C and below. The HfSiO4 phase may be  formed  upon  heating  or  during  cooling.  Formation  of  HfSiO4 occurs when silicon atoms diffuse into the HfO2  crystal interstitially until the solubility limit is reached and has been shown to occur as low as 1100-1400 °C [20, 23].  This can occur  in both Layer  I between the HfO2 crystals  embedded in the SiO2 matrix and Layer  II when the SiO2  fills the HfO2 matrix. The formation of a crystalline solid  from an amorphous glass may help limit oxygen penetra tion and therefore limits oxide scale growth; however,  the  HfSiO4 phase is not continuous and there remains a sub stantial oxidation pathway through the  amorphous SiO2 rich glass. Additionally,  there does not seem to be a cor relation with the degree of oxidation and the presence of  HfSiO2.  The appearance of the tetragonal or cubic HfO2 peaks in the XRD was observed for samples heated to 1700 °C and  above making  it  likely  that  the  phase  transformation  is  monoclinic to tetragonal as expected from the Hf-O phase  diagram [19, 20]. This phase transformation from mono clinic to tetragonal HfO2 results in about 3.5% reduction in  volume [4]. The monoclinic to tetragonal transition may be  linked to the transition from equiaxed to columnar HfO2 grains as reported previously around 1650 °C [4, 24]. The  resulting volume decrease opens additional porosity in the  HfO2  matrix  aiding  the  increased  SiO2  penetration  observed  at  these  temperatures  and making  the  grains  appear more  elongated.  In  pure metal  diborides,  the  transformation from monoclinic to tetragonal is shown to greatly enhance the oxidation of HfB2 at *1700 °C [25,  26]. However,  the existence of SiO2 appears to delay this  effect since the total weight gain is less than 3% of the starting sample weight up to 1800 °C and doubles to 6% at 1900 °C.  A change in morphology may also be caused by a flow  of borosilicate glass  containing HfO2. Studies  involving  ZrB2-SiC [27, 28] discuss the convection of ZrO2-B2O3-  SiO2  liquid driven to the  surface by the volume change  associated with oxidation The ZrO2 then precipitates out of  the liquid as it reaches the vicinity of the surface when the  B2O3 evaporates. The metal oxide (MeO2) phase is pre dicted to precipitate due  to the  fact  that  its  solubility is  higher  in a borosilicate glass than in a pure silicate glass  [27]. The phase diagrams of HfO2-SiO2  [19] and ZrO2-  SiO2 [29] show that about 2 mol% ZrO2 or HfO2 may be  dissolved in the pure SiO2 liquid at the eutectic point (1687 °C for ZrO2-SiO2 and around 1450 °C for HfO2-  SiO2), while  the  amount of dissolvable MeO2 2000 °C approximately  increases  with  temperature. At  40 mol%  HfO2 or 11 mol% ZrO2 would be soluble in SiO2 before  precipitating  as  a  tetragonal  crystal. Upon  cooling,  the  excess HfO2 would precipitate in the SiO2 layer. The flow  of MeO2-B2O3-SiO2  liquid  is  predicted  to  leave  an  underlying  columnar MeO2  layer  and  smaller  equiaxed  oxide particles near or in the SiO2 layer. At 1400 and 1500 °C less than 25 distinct HfO2 grains are observed in  the SiO2 layer (Layer I) in the cross section of the samples.  These grains are identified by their contrast difference with SiO2 and confirmed with EDS. At 1600 °C the population of 1-2-lm HfO2 grains within the SiO2 2000 °C  layer  increases  (Fig. 3)  and  at  1900  and  the  grains  exist  throughout the SiO2 scale. A systematic oxidation study of ZrB2-SiC between 1600 and 1900 °C with hold times of  1 h [28] showed similar oxide scales as observed in this study. At 1600 °C, a SiO2 layer covers a ZrO2 layer whose  pores are filled with SiO2. A SiC-depleted region is not observed until 1700 °C, while oriented growth of ZrO2 is observed at 1800 °C. At 1800 °C,  the SiC-depleted region  is said to contain ‘‘minor’’ amounts of Si, but  there is no  discussion as to whether this is in the form of SiO2, SiC, or 1900 °C,  something  else. At  a  SiC-depleted  layer  is  observed in ZrB2  similar  to Layer IV found for 1800 °C and  the oxi dation  of HfB2-SiC at  higher.  Since  the  monoclinic to tetragonal transformation ZrO2 is between 1000 and 1150 °C,  temperature  of  the structural change  of ZrO2 from equiaxed to columnar between 1700 and 1800 °C would not be expected to be caused by a phase  transformation. Therefore,  the  columnar growth habit  in  J Mater Sci  (2009) 44:5673-5681  5679  123  \\x0c', 'both HfB2-SiC and ZrB2-SiC are from the interaction of  the MeO2 with the borosilicate glass and its convection. 2000 °C in the HfB2-SiC system  Between  1400 and  there are four competing factors contributing to oxidation:  (1) the formation of HfO2, (2) the formation and stability of  the SiO2  scale,  (3)  the presence of HfSiO4,  and (4)  the  transition in morphology and increase in porosity caused  by  a  phase  transformation  and/or  convection  of HfO2-  B2O3-SiO2  liquid.  From the  observations,  it  seems  as  though  the  stability  of  the  SiO2  layer  plays  the most  important  role  in oxidation protection with the  apparent  opening of porosity playing a supporting role. Methods that  have  been  proposed  to  increase  oxidation  resistance  include  engineering the glassy phase by either  inducing  phase separations by the addition of Group IV-VI  transi tion metals  [11] or by the addition of different combina tions  of  Na2O,  Al2O3,  and/or  B2O3  to  incorporate  additional HfO2 [9, 30]. However, decreasing the solubility  of HfO2 may aid in slowing the transport of HfO2 from the  porous  oxide  scale  to  the  SiO2  amorphous  scale  and  thereby  suppresses  the  formation  of  additional  porosity.  The incorporation of TaSi2 has been shown to increase oxidation resistance of ZrB2-SiC ceramics up to 1600 °C,  but has conflicting results at higher  temperatures  [6, 31].  Additionally, stabilizing agents may be alloyed within the  UHTC to delay or halt  the monoclinic to tetragonal phase  transformations of  the HfO2.  Comparison between testing variations  There is no significant difference between samples heated  in  the MoSi2  and ZrO2  furnaces. Comparisons  of  the  weight gain and morphology (not shown) of the samples heated to 1600 °C in both furnaces show the same results.  In  arc  jet  testing  one might  expect  a  catalytic  atom  recombination effect due to the presence of dissociated O  and N present  in the heating plasma of the arc jet (and also  at  the  leading  edges  in  hypersonic  flight);  however,  recombination effects at  the surface of ZrB2 and HfB2 and  particularly their oxides have been shown to be negligible  [32, 33]. Additionally,  the stagnation pressure inside the  arc  jet may  be  less  than  1 atm (in  this  test  stagnation  pressure was 0.1 atm) which can influence  the  active  to  passive transitions in the oxidation of  the samples. This is  especially  true with  the  oxidation  of SiC. When  active  oxidation of SiC occurs  the protective SiO2  layer can be  removed  as  gaseous  SiO.  This  active  oxidation may  account  for  the  very  thin  (0-0.8 lm)  outer  SiO2  layer  observed on the arc jet heated sample when compared to the *0.5-5-lm SiO2  layer  formed on the furnace-heated  sample. This  lack of outer SiO2 may lead to accelerated  oxidation and the longer oxide scale observed for the arc jet heated sample (*18 lm) compared to the sample that was  furnace heated for 4 min (*7 lm). Rezaie et al. [34] have shown that ZrB2-SiC heated to 1500 °C in a lower oxygen  partial pressure has a thinner outer SiO2 layer and a thicker  overall oxide  scale when compared to the  same  sample  heated in air. Some SiO2 may also be lost due to the shear  forces  of  the  plasma.  These  effects  would  become  increasing important with increases  in temperature as  the  SiO2 becomes more fluid. This effect does not change the morphology and chemistry of the oxide scale near 1500 °C  as the arc jet heated sample and the furnace-heated samples  held at 4 min and 1 h all consists of a SiO2 outer  layer  (Layer I) and a porous HfO2 layer with Si-O-C inclusions  (Layer II).  Arc jet tests published by Gasch et al. [11] with a steady-state temperature between 1690 and 1750 °C for a  total exposure time of 20 min can be used as a comparison for the high-temperature tests ([1600 °C) in this study. As in our 1600 and 1700 °C furnace-heated samples, the oxide  scale was described to consist mainly of a porous HfO2  layer. They also report an HfB2  layer with SiC removed  (there is no mention of Si-O-C inclusions in the descrip tion of the arc jet sample). The total oxide scale thickness  of the arc jet heated sample was 72 lm when compared to the 39-lm total oxide scale observed in the 1700 °C fur nace-heated sample of this study that was held at time. From our 1500 °C data  temper ature for a longer  and the  literature data, one can conclude that  the furnace tests can  reproduce the basic microstructures of the oxidized HfB2-  SiC found in arc jet  tests, but  the total  scale thickness  is  limited  by  a  thicker  protective SiO2  outer  layer  in  the  furnace tests.  Conclusion  Samples were heated in air  from 1400 to 2000 °C using a  MoSi2 and ZrO2  furnace. These  samples were  shown to  develop four layers over this temperature range. Below 1800 °C, the oxide scale consists of three layers: Layer I, a  protective  SiO2; Layer  II,  SiO2  infiltrating  into  porous  HfO2; and Layer III, predominantly porous HfO2. At 1800 °C and above, the oxide layer consists of three layers:  Layer  I, a protective SiO2; Layer  II, SiO2 infiltrating into  porous HfO2; and Layer IV, predominately HfB2. Regions  of the predominately HfO2 and HfB2 layers both contained  inclusions of Si-O-C indicating partially oxidized SiC. Up to 2000 °C,  the critical  factors influencing oxidation were  shown to be the flow of SiO2 glass and an apparent increase  in porosity of the HfO2 oxide scale. The outer oxide scale  must be protective against oxygen penetration and as shown  by the  comparison of  furnace-heated samples  to arc  jet  heated samples must be  resistant  to active oxidation.  In  order to increase the working times at elevated temperatures  5680  J Mater Sci  (2009) 44:5673-5681  123  \\x0c', '([1800 °C)  it  is necessary to develop a more robust pro tective glassy scale through the use of additives with SiO2 to  increase the viscosity or melting temperature or to stabilize  the HfO2 crystalline phase.  Acknowledgements  This study was supported in part by the United  States Air Force Contract # FA8650-04-D-5233 with project manager  Michael Cinibulk. We acknowledge NASA-SCAP for  their critical  financial support of the arc jet operational capability at Ames. And we  would like to thank Sylvia Johnson and Matthew Gasch at NASA  Ames for  their assistance in performing the arc jet  testing.  References  1. Monteverde F, Scatteia L (2007) J Am Ceram Soc 90(4):130  2. Fenter JR (1971) SAMPE Quart 2(3):1  3. Fahrenholtz WG, Hilmas GE, Talmy  IG, Zaykoski A (2007)  J Am Ceram Soc 90(5):1347  4. Kaufman L, Clougherty EV, Berkowitz-Mattuck JB (1967) Trans  Metall Soc AIME 239:458  5. Hinze  JW, Tripp WC, Graham HC (1975)  J Electrochem Soc  122(9):1249  6. Opila E, Levine S, Lorincz J  (2004)  J Mater Sci 39:5969. doi:  10.1023/B:JMSC.0000041693.32531.d1  7. Lide DR (1996) CRC handbook of chemistry and physics, 77th  edn. CRC Press Inc, Boca Raton (Sect. 4)  8. Clougherty EV, Pober RL, Kaufman L (1968) Trans Metall Soc  AIME 242:1077  9. Lespade P, Richet N, Goursat P (2007) Acta Astronautica 60:858  10. Monteverde F (2005) Corr Sci 47:2020  11. 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Marschall  J, Chamberlain A, Crunkleton D, Rogers B (2004)  J Spacecr Rocket 41(4):576  33. Scatteia L, Borrelli R, Cosentino G, Beˆ che E, Sans  J-L, Balat Pichelin M (2006) J Spacecr Rocket 43(5):1004  34. Rezaie A, Fahrenholtz WG, Hilmas GE (2006) J Am Ceram Soc  89(10):3240  J Mater Sci  (2009) 44:5673-5681  5681  123  \\x0c']"
},{
  "_id": 193,
  "PDF": "Oxidation resistance of HfB2–SiC composites for protection ofcarbon-based materials.pdf",
  "Text": "['Acta Astronautica 60 (2007) 858 - 864  www.elsevier.com/locate/actaastro  Oxidation resistance of HfB2-SiC composites for protection of carbon-based materials  P. Lespadea , ∗  , N. Richetb , P. Goursatb  a EADS Space Transportation-33165 St Médard en Jalles, France b SPCTS UMR 6638 Université de Limoges-87060 Limoges, France  Received 15 February 2006;  received in revised form 28 July 2006; accepted 5 November 2006  Available online 12 January 2007  Abstract     flowing O2 /He mixtures (Ptot = 1 bar, PO2 = 1-200 mbar). Two overlapping domains of oxidation were identified corresponding, The oxidation behaviour of HfB2 -SiC monoliths and coatings has been studied in the 20-1700 C temperature range in respectively, to the reaction of HfB2 and an intermediate phase containing Hf/Si/B/C in the 600-800 C temperature range and SiC nanoparticles in the 800-1000 C interval. A protective borosilicate glass is quickly formed at low temperature. The good oxidation resistance up to 1700 C and under low oxygen partial pressure (10 mbar) is discussed according to the degradation mechanisms involving compositional and microstructural changes and compared to previous results obtained on hot pressed HfB2 -SiC composites. The specific microstructure of the materials studied in this work is of great importance for the composition of the protective glassy phase and leads to the formation of a refractory glass containing hafnium. © 2006 Elsevier Ltd. All rights reserved.           Keywords: Carbon/carbon composites; Graphite; Coating; Oxidation; Sintering  1.  Introduction  Potential ceramics that can be used at very high temperatures are carbides, silicides, borides and nitrides [1-7]. Among these candidates, literature data show that carbides and borides which have high melting temperatures, possess the required thermomechanical properties for thermostructural applications [8,9]. However, in oxidizing environments the degradation mechanisms which correspond to a combined inward diffusion of oxygen and the release of gaseous species (CO, CO2 , B2O3 ) restrict their applications. Therefore, the  ∗  Corresponding author.  E-mail address: pierre.lespade@space.eads.net  (P. Lespade).  0094-5765/$ see front matter © 2006 Elsevier Ltd. All  rights reserved.  doi:10.1016/j.actaastro.2006.11.007  oxidation resistance is mainly influenced by the microstructure and the plasticity of the oxide scale. To improve the oxidation behaviour of HfB2 or ZrB2 , it was suggested to incorporate appropriate additives containing silicon [10,11]. HfB2 -SiC and ZrB2 -SiC composites were elaborated by hot pressing and their reactivity with oxygen was evaluated up to 2000 C [5,10,12,13]. Due to the selective oxidation of HfB2 grains at low temperature and the escape of boron containing species before the oxidation of SiC, a porous HfO2 scale is formed with an external and protective silica layer. Thus, the use of these composites is limited by the stability of silica at high temperature or under low oxygen partial pressure. The aim of this work is to study the oxidation resistance of HfB2 -SiC composites that exhibit a specific microstructure leading to a     \\x0c', 'P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  859  complex glassy phase at high temperature in order protect carbon-based materials.  to  2. Experimental procedure  2.1. Materials  The oxidation behaviour of HfB2 -SiC materials was studied for monoliths (200-400 ♯m thickness) or for coatings (100-200 ♯m thickness) on carbon-based materials (graphite, C fibres/C matrix or C fibres/SiC matrix composites). Three processes [14] were developed for monolith or coatings, leading to materials with various microstructures and compositions as summarized in Table 1. The pore size distribution measurements were conducted with a mercury porosimeter (Micromeritics Autopore 9215). Table 1 shows that the pore size diameter is higher for material B, but the total porosity is less than 10-12% for all the samples. The crystalline phases were identified by X-ray diffraction (Siemens D5000) and transmission electron microscopy (TEM JEOL 2010). Composites consist of HfB2 crystals 2-5 ♯m embedded with with a grain size of SiC nanoparticles. For material C, the presence of HfN1−x B2x solid solution is detected by X-ray diffraction and Auger Electron Spectroscopy (Thermo VG Scientific Microlabs 310F). The local composition was also determined by AES for all the samples. The concentration profiles of the elements at the interface of HfB2 crystals and the SiC matrix evidenced the presence of a Hf/B/Si/C phase and compositional gradients on 0.4-0.7 ♯m distance. After grinding the samples, the chemical analysis of the powders was performed (Centre d’analyse CNRS Vernaison France). Due to the presence of a nonstoechiometric Hf/B/Si/C phase around HfB2 grains it is very difficult to determine the mineral composition of materials A-C.  2.2. Oxidation behaviour  Oxidation kinetics were performed using a thermogravimetric analysis equipment (Setaram). The mass  Table 1  Materials main properties  Materials  Mean pore diameter  (♯m)  Crystalline phases (XRD)  A  B  C  0.11  0.30  0.10  traces  HfB2 , SiC, Si HfB2 , SiC HfB2 , SiC, HfN1−x B2x     changes were followed with a linear increase of temperature (5 /min) or under isothermal conditions with constant flow of a O2 /He mixture (50 ml/min). For isothermal runs the samples are heated under pure argon until the soaking temperature is reached. Then the oxidizing gas mixture is introduced with a flow of 50 ml/min. Various oxygen partial pressures were used by decreasing the oxygen content in the O2 /He mixture but the total pressure was kept always at 1 bar. After oxidation, the samples were analysed using X-ray diffraction, scanning electron microscopy (SEM JEOL SM 35) and microprobe analysis (Cameca SX100) for the elemental distribution and compositional profiles in the oxide scale. In order to compare our data with literature results, we have plotted ♅w/S (♅w : weight versus temperature or time change, S: geometrical surface of the sample).  3. Results  3.1.  Influence of  temperature           The oxidation behaviour in a O2 /He mixture was studied up to 1700 C. The weight changes for the different materials heated with a linear increase of temperature are reported in Fig. 1. As evidenced by the derived curves the reaction with oxygen starts above 500 C and two overlapping domains of oxidation are detected. In the first domain between 600 and 800 C, HfB2 and the intergranular phase Hf/Si/B/C are oxidized as shown by the analysis of the reaction products (see 3.3). At higher temperatures (800-1000 C) the oxidation of SiC nanoparticles occurs and the overall rate of reaction slows down quickly. One notices that the weight gain is higher for material B due to the size of the porosity. For isothermal conditions an example of the weight gain curves is plotted in Figs. 2 and 3. The shape of the curves with a decrease of the rate of the reaction versus time, whatever the temperature, is in favour of the formation of a protective oxide scale. For all the temperatures studied the advancement of the oxidation is higher for materials B and C but remains small. Due to the different nature of the crystalline phases and the pore size distribution, a monotonous increase of the maximum weight gain with temperature is not observed. This behaviour could be due to the release of boron containing species above 1100 C. If the data are corrected to take into account the escape of gaseous species, for material C oxidized 10 h at 800 C the weight change is higher than at 1600 C, which              \\x0c', '860  P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  0  200  400  600  800  1000  1200  1400  1600  0  1  2  3  4  5  6  7  8   material A   material B   material C  Temperature (°C)  0  200  400  600  800  1000  1200  1400  1600  0  1  2  3  4   material A   material B   material C  d  i  f f  (  Δ  w  /  S  )  Δ  w  /  S  (  m  g  /  c  m  2  )  Temperature (°C)  Fig. 1. Linear mixture with PO2 = 200 mbar : increase of temperature on monoliths under O2 /He (a) mass variation and (b) derivative curves showing two overlapping domains.  confirms the preferential oxidation of HfB2 Hf/B/Si/C phase at low temperature. Concerning the small increase of the weight gain at 1600 C, it can be attributed to a balance between the release of volatile species and the oxide formation or to a very slow oxidation rate. The thicknesses of the substrates and oxide scales were measured after oxidation during 3, 5, and 10 h under 20% O2 /He mixture at 1600 C in order to determine the recession rate. The oxide scale grows slowly from 20 to 24 ♯m whatever the thickness of the starting sample. Moreover, as shown later the oxide scale is stable up to 1700 C under 200 mbar of oxygen (Fig. 6) which confirms the slow oxidation process.  and the           0  5  10  0  1  2  3  4  5  6  7  8  9  10  11  12  13  14  15  800°C 1000°C 1450°C 1600°C  Δ Δ  w  /  S  (  m  g  /  c  m  2  )  1  2  3  4  TIME (h)  6  7  8  9  PO2 = 200 mbar material C. Fig. 2. Isothermal oxidation of monoliths under O2 /He mixture with  0  5  6  0  1  2  3  4  5  6  7  8  9  10  11  12  13  14  15  800°C  1000°C  1450°C  1600°C  Δ  w  /  S  (  m  g  /  c  m  2  )  1  2  3  4  TIME (h)  7  8  9  10  PO2 = 50 mbar material C. Fig. 3. Isothermal oxidation of monoliths under O2 /He mixture with  3.2.  Influence of oxygen partial pressure  In order to study the thermal stability of the oxide scale versus oxygen pressure, the oxidation behaviour of the material A was investigated in the 1-200 mbar range. At a very low oxygen partial pressure, a release of SiO from the oxide scale, a phenomenon similar to the transition between passive and active oxidation of silicon-based-materials, could be observed [15,16]. The transition temperature corresponding to a weight loss was determined by TGA during a linear increase   / min) of temperature (1 for 1, 5, 10 and 200 mbar (PO2 + PHe = 1 bar). As of oxygen the transition              \\x0c', 'P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  861  Temperature (°C)  1100  1200  1300  1400  1600  1800  Keys  PASSIVE  Gulbr. (A)  Vaughn  Balat  Gulbr.(C)  this study  Hinze  ACTIVE  monoliths HfB2-SiC SiC-β  powder  ) r  a b  m  (  2  O  P  100  10  1  0.1  0.01  1E-3  -7.5  -7.0  -6.5  -6.0  -5.5  -5.0  -4.5  -10000/T (K-1)      Material C was also oxidized under isothermal conditions. The behaviour observed at 1600 C (Fig. 5) confirms the results obtained with linear increase of temperature. The weight loss at 1 mbar corresponding to a release of silicon and oxygen from the oxide scale is due to the decomposition of the silicate (see 3.3). However, the degradation kinetics of the reaction products under low oxygen pressure are different from active oxidation of silicon or silicon carbide. A linear region with catastrophic degradation of our HfB2-SiC material is not observed. The porosity formed with the escape of the volatile species from the oxide scale is partially sealed by the vitreous phase. The oxidation rate decreases progressively with the advancement of the reaction.  Fig. 4. Oxygen partial pressure versus  temperature  for  the  active  to passive oxidation 300 cm3 /min).  transition  of SiC-based materials  (flow rate  3.3. Reaction products  15  10  5  0  -5  )  2  m  c  /  g  m  (  S  /  w  Δ  -10  -15  -20  10 mbar  5 mbar  1 mbar  0  1  2  3  4  5  6  7  8  9  10  TEMPS (h)  Fig. 5. Oxidation of monolith (material C) at 1600 C under various oxygen partial pressures and a flow rate of 300 cm3 /min.     temperature depends strongly on the experimental conditions [15,16] a silicon carbide powder (♢-SiC 99.8%) was oxidized with the same experimental conditions. Our results are compared with previous works on silicon-based materials and summarized in Fig. 4. HfB2-SiC composites exhibit a larger passive oxidation domain compared to silicon carbide powder. This graph shows that the stability of the oxide layer formed on HfB2-SiC materials is higher than silica at every oxygen partial pressure tested. The transition is reached under PO2 = 10 mbar not at 1700 C, which indicates that such a protection is still efficient in these conditions.        X-ray diffraction patterns of the surface of oxidized samples evidence the presence of HfO2 (ASTM 340104) and a vitreous phase for all the temperature range studied. A dense and pore free oxide layer is formed on the substrates up to 1700 C as shown by SEM (Fig. 6). In order to determine the composition and the elemental distribution, microanalysis was performed on sample cross sections. After an oxidation at 800 C for 10 h, a thin oxide layer containing Hf/B/Si/O is formed. Traces of oxygen are observed in the bulk of the HfB2-SiC samples due to the presence of a limited porosity which agrees with the weight gain curves (Fig. 2). The thickness of the oxide scale increases slowly with temperature but porosity is never observed. At 1450 C the vitreous phase contains mainly silicon, oxygen and a small content of boron. After 10 h at 1600 C, under 200 mbar of oxygen, a protective scale is formed with HfO2 crystals and vitreous phase (Fig. 7). Concentration profiles in this glassy phase show that a small quantity of HfO2 is dissolved. Moreover, a boron gradient from the internal interface to the external interface is also evidenced which indicates a release of boron species from the scale. Microanalysis of the vitreous phase was also performed from the internal interface to the external interface on oxidized samples at 1700 C for 10 h under 200 mbar of oxygen. The content of dissolved HfO2 increases with the oxidation temperature but the glass contains also boron in an area close to the internal interface (Table 2).                  \\x0c', '862  P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  Table 2  Microprobe analysis: chemical composition of the glassy phase, after     1600  C/10 h under PO2 200 mbar  Measure  Atomic concentration (%)  1  2  3  O  67.4  71.2  69.6  Si  22.1  28.3  29.8  B  8.5  0  0  Hf  2.0  0.5  0.6  4. Discussion  In an oxidizing environment HfB2 reacts with oxygen according to the following reaction: HfB2 (s) + 5/2O2 (g) → HfO2 (s) + B2O3 (l).  As mentioned in the literature, the diffusion regime of oxygen through liquid boria at low temperature changes in a paralinear regime at high temperature due to the volatilization of boron containing species [2,3,9]. For silicon carbide two oxidation mechanisms are observed according to temperature and oxygen pressure. A passive oxidation occurs for high oxygen pressure: SiC(s) + 3/2O2 (g) → SiO2 (s) + CO(g).  While an active oxidation for high temperature and low oxygen pressure is observed: SiC(s) + O2 (g) → SiO(g) + CO(g).  section of a coating elaborated on a C/SiC  C/10 h under 20%O2 /He mixture and a flow  Fig. 6. Polished cross  composite, after 1700 rate of 50 cm3 /min.     For HfB2-SiC composites these two behaviours can be modified by interactions between the reaction products. The thermal stability and the boiling temperature of the borosilicate are higher than for boria [17]. In order to explain the influence of the microstructure of HfB2 -SiC monoliths or coatings, the oxidation behaviour is compared with previous results on hotpressed materials with nearly the same nominal composition. In the case of hot-pressed HfB2 -SiC composites the grain size is of 3-5 ♯m for SiC and 5-10 ♯m for HfB2 [4,9,10]. When specimens are heated under 250 mbar of oxygen, a selective oxidation of HfB2 starts above 600 C leading to the formation of HfO2 and liquid boria, which losses boron containing species as temperature is raised. At the beginning of SiC oxidation (1100 C), most of the B2O3 is evaporated. SEM micrographs and microprobe analysis show that the oxC is 50 ♯m thick and composed of ide scale at 1400 an amorphous silica layer on the top of a porous HfO2 layer.           Fig.  1600     7. Microprobe  analysis  of  a monolith  cross  section  after  C/10 h, 20%O2 /He mixture (material A).  \\x0c', 'P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  863  0  4  10  1450°C 1600°C  (  Δ  w  /  S  ))  2  (  m  g  2  .  c  m   4  )  2.0  1.5  1.0  0.5  0.0  2  TIME (h)  6  8  Fig. 8. Parabolic transformation of mass variations versus time.  On the contrary, for this work HfB2 -SiC monoliths or coating a microstructure with fine SiC grains is developed. The oxidation of HfB2 crystals and the Hf/B/Si/C phase induces the formation of HfO2 and a borosilicate glass from 600 to 700 C. The increase of silica content when SiC nanoparticles are oxidized (T > 800 C) improves the stability of the borosilicate glass. However, its viscosity is low enough to seal progressively the porosity of the substrate and to form a dense protective layer. Pores or bubbles are never observed in the oxide scale which is thinner (25-30 ♯m after 10 h at 1600 C under 200 mbar of O2 /He mixture) than for hot-pressed samples. The release of boron containing species is shifted to higher temperatures as evidenced by microprobe analysis and the temperature interval between the oxidation of HfB2 and SiC mentioned by Hinze [5] is not observed for our monoliths and coatings. Moreover, HfO2 dissolution in the borosilicate glass is detected above 1600 C. Concentration profiles show that hafnium content is higher in a domain close to the internal interface. These analyses agree with results published in previous paper [18]. The use of Na2O, Al2O3 or B2O3 facilitates the incorporation of HfO2 in silicate glasses. Taking into account the shape of the weight gain curves at 1450 and 1600 C and the morphology of the oxidation products we have plotted (♅w/S )2 versus time (Fig. 8). For the different materials a parabolic behaviour is observed after an initial transient period. The parabolic rate constants calculated from our experiments are compared to the results of previous works on hot-pressed samples having nearly the same composition (Fig. 9). The diffusion rate constants for monoliths or coatings are smaller than for hot-pressed composites. Moreover, one notices an influence of oxygen partial pressure                 1600  1500  1400  1200  1300  after oxidation under 333mbar(10)  after oxidation under 200mbar  after oxidation under 50mbar  5.2  5.4  5.6  5.8  6.0  6.2  -7.4  -7.2  -7.0  -6.8  -6.6  -6.4  -6.2  -6.0  -5.8  -5.6  -5.4  -5.2  -5.0  after oxidation under 200mbar  l  g o  K  p  (  g  2  .  c  m   4  .  h   1  )  104/T (K-1)  HIP monolith  material A  material C  material C   Temperature (°C)  Fig. 9. Arrhenius plot of parabolic rate constants—green triangles  are for  literature data (10) on HIP monolith.  on the parabolic rate constant as in the case of oxidation of silicon-based compounds. This influence is in favour of a limiting step corresponding to oxygen diffusion to the internal interface. The better oxidation resistance of materials A-C in comparison with hot-pressed composites above 1450 C can be explained by the specific microstructure of the materials. The presence of a Hf/B/Si/C phase around HfB2 grains induces the formation of a borosilicate more stable than boria and a dissolution of the HfO2 which grows above 1450 C. The increase of the viscosity of this glass compared to silica could explain the slower rate of oxygen migration and the very good oxidation resistance of the HfB2 -SiC composites up to 1700 C under 200 mbar of oxygen and down to 10 mbar at least. For low oxygen pressures and above 1600 C a release of silicon and oxygen occurs. But the reaction with oxygen is limited to the HfB2 -SiC coating. Microprobe analysis evidences a glassy phase in the porosity of the oxide scale and around unoxidized HfB2 grains. After a transient period, compositional changes become very slow and the carbon base substrate remains protected against oxidation for many hours.              5. Conclusion  A good oxidation resistance from 600 to 1700 C down to 10 mbar oxygen partial pressure has been           \\x0c', '864  P. Lespade et al.  / Acta Astronautica 60 (2007) 858 - 864  obtained on HfB2 -SiC composites. Two overlapping domains of oxidation were identified. In the 600-800 C temperature range, HfB2 and the Hf/Si/B/C phase are oxidized leading to a protective borosilicate scale. In the 800-1000 C temperature interval, SiC nanoparticles react with oxygen and the enrichment of the glassy phase with silica at low temperature slows the release of boron containing species. This borosilicate remains up to 1700 C at the internal interface between the oxide scale and the unoxidized material. HfO2 dissolution in the borosilicate has been evidenced and is facilitated by the presence of boron. This dissolution of hafnium in the glass could explain the good stability of the glassy phase up to 1700 C under a 1%O2 /He mixture. Transition between active and passive oxidation behaviour has been characterized as observed on silicon carbide-based materials but for HfB2 -SiC composites the passive oxidation domain is wider.              References  [1] C.B.  Bargeron,  R.C.  Benson,  R.W. Newman, A.N.  Jette,  Oxidation  mechanisms  of  hafnium  carbide  and  hafnium  diboride  in  the  temperature  range  1400  to  2100     C,  Johns  Hopkins APL Technical digest 14 (1)  (1993) 29-35.  [2] J.B.  Berkowitz-Mattuck,  High  temperature  oxidation.  III  Zirconium  and  Hafnium  diborides,  Journal  of  the  Electrochemical Society (1966) 908-914.  [3] J.B.  Berkowitz-Mattuck,  High  temperature  oxidation.  IV  zirconium and hafnium carbides, Journal of the Electrochemical  Society (1966) 1030-1033.  [4] A. Derre, M. Ducarroir,  F. Teyssandier, Réflexions  sur  les  possibilités  d’amélioration  de  la  résistance  à  l’oxydation  des matériaux  carbonés,  Revue  Internationale Des Hautes  Temperatures et des Refractaires 29 (1994) 11-35.  [5] J.W. Hinze, W.C. a HfB2 + 20v/o SiC composite, Tripp, The high temperature Electrochemical Society 122 (9) (1975) 1249-1254.  oxidation  behavior  of  Journal  of  the  [6] A.G. Metcalfe, N.B. Elsner, D.T. Allen, E. Wuchina, M. Opeka,  E. Opila, Oxidation of hafnium diboride, The Electrochemical  Society Proceedings 99-38 (2000) 489-501.  [7] Y. Xu,  L. Cheng,  L.  Zhang, H. Ying, Oxidation  behavior  and mechanical properties of C/SiC composites with Si-MoSi2  oxidation protection coating,  Journal of Materials Science 34  (1999) 6009-6014.  [8] L. Kaufman, Borides  composites—a new generation of nose  cap  and  leading  edge materials  for  reusable  lifting  re-entry  systems, AIAA Advanced Space Transportation Meeting, Cocoa  Beach, FL, 1970.  [9] S.R. Levine, E.J. Opila, Evaluation of ultra-high temperature  ceramics  for  a  eropropulsion  use,  Journal  of  the European  Ceramic Society 22 (14-15)  (2002) 2757-2767.  [10] E. Clougherty, D. Kalish, E. Peters, Research and development  of  refractory  oxidation  resistant  diborides. Technical Report  AFML-TR-68-190, 1968.  [11] W.C. Tripp, H.H. Davis, H.C. Graham, Effect of an SiC addition  on the oxidation of ZrB2 , Ceramic Bulletin 52 (1973) 612-616. [12] E. Clougherty, R. Hill, W. Rhodes, E. Peters, Research and  development of  refractory oxidation-resistant diborides. Part  II  Processing and characterisation. Technical Report AFML-TR 68-100, 1970.  [13] L.  Kaufman,  E.V.  Clougherty,  Oxidation  characteristics  of  hafnium and  zirconium diboride,  Transactions  of  the  Metallurgical Society of AIME 239 (1967) 458-466.  [14] P. Lespade, N. Richet, P. Goursat, Procédé de fabrication d’un  matériau  réfractaire  revêtement  protecteur  susceptible  d’être  obtenu par ce procédé et  leur utilisation. FRA0209677.  [15] M.J.H. Balat, Determination of  the active to passive transition  in the oxidation of silicon carbide in standard and microwave excited air, Journal of the European Ceramic Society 16 (1996)  55-62.  [16] W.L. Vaughn, Active-to-passive  transition in the oxidation of  silicon carbide and silicon nitride in air, Journal of the American  Ceramic Society 73 (6)  (1990) 1540-1543.  [17] D.W. McKee,  Oxidation  protection  of  carbon  materials,  Chemistry and Physics of Carbon 23 (1991) 173-232.  [18] L.L. Davis, G.G. Li, The effects of Na2O, Al2O3 and B2O3 on HfO2 solubility in borosilicate glass, Materials Research Society Symposium Proceedings 556 (1999) 313-320.  Further reading  [19] J.B. Davis, D.B. Marshall, K.S. Oka, R.M. Housley, P.E.D.  Morgan, Ceramic  composites  for  thermal protection systems.  Composite Part A 30 (1999) 483-488.  [20] W.C. Tripp, Thermogravimetric study of the oxidation of ZrB2 in temperature range of 800 to 1500 C, Journal of the        Electrochemical Society 118 (7)  (1971) 1195-1199.  \\x0c']"
},{
  "_id": 194,
  "PDF": "Oxidation resistance of tantalum carbide-hafnium carbide solid solutions under the extreme conditions of a plasma jet.pdf",
  "Text": "['Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Oxidation resistance of tantalum carbide-hafnium carbide solid solutions  under the extreme conditions of a plasma jet  Cheng Zhang, Benjamin Boesl, Arvind Agarwal  ⁎  Plasma Forming Laboratory, Department of Mechanical and Materials Engineering, Florida International University, Miami, FL, USA  A R T I C L E  I N F O  Keywords:  Carbides  Oxidation resistance  Solid solution  Spark plasma sintering  Ultra-high temperature ceramics  A B S T R A C T  The oxidation behaviors of  tantalum carbide (TaC)hafnium carbide (HfC) solid solutions with five different  compositions,  pure HfC, HfC-20 vol% TaC (T20H80), HfC 50 vol% TaC (T50H50), HfC 80 vol% TaC  (T80H20), and pure TaC have been investigated by exposing to a plasma torch which has a temperature of  approximately 2800 °C with a gas flow speed greater than 300 m/s for 60 s, 180 s, and 300 s, respectively. The  solid solution samples  showed significantly  improved oxidation resistance  compared to the pure  carbide  samples, and the T50H50 samples exhibited the best oxidation resistance of all samples. The thickness of the  oxide scales in T50H50 was reduced more than 90% compared to the pure TaC samples, and more than 85%  compared to the pure HfC samples after 300 s oxidation tests. A new Ta2Hf6O17 phase was responsible for the improved oxidation performance exhibited by solid solutions. The oxide scale constitutes of a  found to be  scaffold-like structure consisting of HfO2 and Ta2Hf6O17 filled with Ta2O5 which was beneficial to the oxidation resistance by limiting the availability of oxygen.  1.  Introduction  Hafnium carbide  (HfC)  and  tantalum carbide  (TaC)  both are [1-3].  considered  ultrahigh  temperature  ceramics  (UHTCs)  Applications  of  these materials  include  rocket nozzles  and leading  edges  for hypersonic  vehicles due  to  their  extremely high melting  points  (~ 3900 °C)  [4]. Even with  such  high melting  points,  the  adoption  of  TaC and HfC have  still  been  obstructed  due  to  the  unsatisfying  oxidation  resistance  of  these  two  carbides,  and  the  oxidation mechanisms are yet to be fully understood.  Early studies have been carried out  to understand the oxidation  behaviors of TaC and HfC, respectively. In general, HfC is considered a  superior  oxidation  resistance material  due  to  a  three-layer  oxide  structure observed in previous isothermal oxidation experiments. The  oxidized structure  consisted of an outer  layer  that  is  fully oxidized  HfO2, an inner layer of un-oxidized carbides, and a partially oxidized, oxy-carbide layer sandwiched in between [5]. The oxy-carbide layer can  act as oxygen diffusion barrier and protect the underlying carbides. The  formed HfO2 has a high melting point around 2900 °C, and it has a very low vapor pressure (10−4 Torr at 2500 °C [7]), which makes it one  of  the least volatile oxides [6]. However, during the oxidation of HfC,  gaseous products like CO and CO2 leave a porous HfO2 oxide scale, which can be detrimental to the mechanical integrity [8].  One solution to enhance the integrity of a porous exterior oxide  layer  is  to seal  the  cracks  and pores with a  liquid phase material,  resulting in a solid scaffold and liquid phase structure. An example of structure can be seen during the oxidation of HfB2-SiC [9-11]. Previous results have demonstrated that HfO2 provides a stable solid scaffold for the mechanical rigidity of the resultant oxide layers to  this  withstand the harsh environment and borosilicate glass melts and seals  the cracks in the HfO2, mechanisms are demonstrated in this system, this specific borosilicate  leading to a fully dense oxide layer. While the  glass begins to evaporate around 1400 °C, limiting the peak application  temperature of  the HfB2-SiC system. Mimicking this higher temperature materials could lead to improved oxidation resis structure with  tant UHTCs with enhanced mechanical rigidity. As a scaffold material,  the enhanced mechanical properties of the UHTC carbide systems have  the potential  to form a stronger scaffold-liquid structure. Candidates  for the liquid sealing phase should have a melt temperature below the  application temperature (~ 2000 °C) to ensure liquid phase formation  and flow but have a boiling point much higher  than the application  temperature  to prevent  the  evaporation (seen in Hf2B-SiC). study, HfC can provide a stable HfO2 solid scaffold, similar to HfB2, while the and the main oxidation product of TaC, Ta2O5 that has a relatively low melting point of 1872 °C, [12] can be the  In this  perfect  candidate to provide liquid phase during the oxidation of HfC while  maintaining stability.  Recently,  few studies on the synthesis of TaC-HfC solid solutions  http://dx.doi.org/10.1016/j.ceramint.2017.07.227  Received 10 July 2017; Received in revised form 30 July 2017; Accepted 31 July 2017  ⁎ Corresponding author.  E-mail address: agarwala@fiu.edu (A. Agarwal).  Ceramics International 43 (2017) 14798-14806  Available online 01 August 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  MARK  \\x0c', \"have emerged [13-15]. In a recent study, we were able to synthesize  compacts  of TaC-HfC solid solutions using  spark plasma  sintering  (SPS) [13]. Five compositions were chosen, pure TaC (PT), TaC-20 vol  % HfC  (T80H20),  TaC-50 vol% HfC  (T50H50),  TaC-80 vol% HfC  (T20H80), and pure HfC (PH)  to explore a range of  compositions.  The formation of a solid solution was achieved in each composition.  The literature on the oxidation behavior of TaC-HfC solid solutions is  very scarce [16,17]. Coutright et al. synthesized HfC and HfC-25 wt%  TaC solid solution by hot pressing. The oxidation behaviors of HfC and  HfC-TaC solid solutions were performed by thermogravimetric method from 1200 to 1530 °C and defocused CO2 laser for 1800-2200 °C. The results showed the microstructure of the oxides of HfC and HfC-TaC  were similar, but with severe cracking in HfC-TaC samples. Ghaffari  et al. studied Ta0.8HfC0.2 solid solution with 24 vol% MoSi2 as sintering aid consolidated using pressureless sintering. The oxidation behavior  was evaluated by oxidized flame at 3000 °C. The results suggested the  formation of double protective oxide layers during the oxidation. None  of  these  two studies provided a clear  verdict on the  effect of  solid  solution formation on oxidation behavior of TaC-HfC based UHTCs.  Moreover,  these  studies  [16,17] were  limited  to  only  a  couple  of  compositions. Our  study  is  the  first  systematic  study  to  evaluate  oxidation  behavior  of  large  spectrum of  TaC-HfC  solid  solution  compositions. Moreover,  oxidation  is  carried  out  in  a  high  speed  plasma  jet with temperatures  exceeding 2000 °C making our  study  unique.  The simulation of the application conditions of UHTCs, a tempera ture  above  2000 °C with  airflow at  sonic  speed,  is  difficult  in  a  laboratory environment. What  is more challenging is how to hold the  samples  under  the  extreme  conditions without  introducing  stress  concentration. Normally  the  specimens are physically held in place  with fixtures, and these fixtures become stress concentrations that lead  to premature cracking during the oxidation testing due to the thermal  expansion of  the specimens. These cracks provide the pathways  for  oxygen to penetrate  and cloud the  true oxidation behaviors of  the  studying subjects. A new fixture has been developed in the current  work and described in Section 2.2. The specimens are held by vacuum  so  that  the  stress  concentration will  be  eliminated. The  oxidation  behavior of TaC-HfC solid solutions samples is evaluated by exposing  to a high-temperature plasma torch that used in plasma spray. The  plasma torch can provide up to 5000 K temperature with airflow with  sonic speed, which is an improved tool to mimic the extreme conditions  in the real world applications. The samples are subjected to the plasma  exposure for 60 s, 180 s, and 300 s. Detailed characterization of  the  oxidized samples have been carried out, and the oxidation mechanisms  have been proposed in the current work.  2. Experimental procedure  2.1. Materials  Oxidation testing was conducted on each composition of the spark  plasma sintered pellets. Each pellet was twenty (20) mm in diameter and approximately 4-5 mm in thickness. The details of  the sintering  conditions can be found in our previous work [13]. The pellets were  ground using 15 µm diamond paper to remove the graphite foils from  the sintering as well as provide relative flat surfaces.  2.2. Oxidation testing  The oxidation tests were  conducted by a plasma flow generated  from a Praxair SG-100 DC plasma gun. The input power was 30 kW  and the exposure time was 60 s, 180 s, and 300 s for all the samples. A  75 mm standoff distance was measured from the plasma  gun and  sample front surface. A significant standoff distance ensured the ample  air  (oxygen)  exposure. Primary  argon gas flowed at 56 slpm,  and  secondary helium gas flowed at 60 slpm to form a plasma. The front  side temperature and gas flow velocity were evaluated by an accura spray in-flight particle diagnostic sensor (Tecnar Automation Ltd., QC,  Canada). The sensor head was mounted 75 mm from the plasma gun,  and  AlO-101(Praxair  Surface  Technologies,  Inc.  Indianapolis,  IN.  USA.) powder was used as spraying powder for the temperature and  velocity measurement. The temperature at  the surface was measured  above 2700 °C with a velocity of 330 m/s. With these simulated testing  conditions,  samples would  experience  both oxidation and ablation  similar to the ones that space vehicles would face upon reentry. The  setup of the newly designed fixture for oxidation tests is schematically  described in Fig. 1a. It consists of a steel tube which is connected to the  vacuum. SPS sample  is held by the vacuum instead of  clamps and  eliminates  stress  concentration  produced  by mechanical  forces.  Moreover,  entire  disk  shaped  sample  is  exposed  to  plasma  (see  Fig. 1b) and not masked by the clamps. A thermocouple was inserted  into the steel  tube to measure the samples' back side temperature. It  was swirled into a coil shape to make sure proper contact throughout  oxidation testing.  2.3. Post-oxidation characterization  The  phases  after the oxidation tests were identified by X-ray (Siemens D-5000) using Cu Kα radiation at a scan  diffraction (XRD)  rate of 2 °/min. The operating voltage and current were set at 40 kV  and 35 mA,  respectively. The relative oxide contents were computed  using the area under the peaks. The oxidized pellets were then cut from  Fig. 1. The oxidation testing setup. (a) A detailed schematic of oxidation testing, (b) snapshot of high-temperature oxidation testing.  C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  14799  \\x0c\", \"the  center  to  expose  the  cross-section. The  exposed  cross-section  surfaces were mounted and polished with the diamond suspension to  0.5 µm finish.  Top  surface  and  cross-section morphologies were  examined by SEM (JEOL 6330F). The front side oxide layer thickness was measured using “Image J”  software based on the cross-section  SEM images. A JEOL-JIB 4500 Multi-Beam Focused Ion Beam (FIB)  was used to expose the solid scaffold structure by milling the cross section of the oxide layers in the solid solution samples.  3. Results and discussion  3.1. Basic information on HfC/TaC solid solutions before oxidation  The basic information on the HfC/TaC solid solutions is provided in  this section and summarized in Table 1 for the sake of completeness of  to the reader. The detailed information about HfC/TaC solid solution  formation and sintering without any aids is provided in our previous  work [13].  All samples have near full densification; the HfC contained samples  have slightly higher densification. The average grain sizes reduce with  the increase of HfC content. The XRD patterns of  the samples show  only the carbide peaks; no secondary peaks were detected [13].  3.2. Oxidation behaviors of TaC, HfC, and TaC-HfC solid solutions  The oxide scales of  the solid solution samples remain intact even  after 300 s exposure. Spallation and delamination are observed on the  pure HfC post-oxidation samples, and especially on the pure TaC post oxidation samples. As a result, only the front side oxide layer thickness  (Fig. 2) is reported in the present work. From Fig. 2,  it is evident that  all  compositions  of  solid solution samples  exhibit  significantly  im proved oxidation resistance as compared to pure carbides. T50H50 has  the thinnest oxide layers for all three durations of oxidation. Its oxide  scale thickness is 1/10 of pure TaC's, and 1/6 of pure HfC's after 300 s  exposure. The oxidation rate can be calculated from the ratio of  the  oxide  layer  thickness  and  testing  duration. After  300 s  tests,  the  oxidation rate for T50H50 is only 0.09 µm/s,  followed by T80H20 is  0.13 µm/s and the T20H80 is 0.34 µm/s. The oxidation rate  is not  available for pure carbides due to the serve delamination during the  oxidation tests. The oxidation rate for solid solutions showed signifi cant  improvement compared to the results reported in the literature  [18]. The ZrB2-SiC system, which oxidation resistance, was reported with an oxidation rate of 0.27 µm/  is  considered  to  have  the  best  s tested by a high-frequency plasma wind tunnel with the surface temperature of 1960 °C and 3-4 g/s gas mass flow. The oxidation rate  is 3 times higher as compared to the T50H50 sample in the present  study.  [18]. To quantify  the oxidation more accurately,  the  carbide  recession rates for 5 samples after 300 s testing were also calculated  and summarized in Table 2. The results suggest  that  the T50H50 has  great ability in preserving the carbide and enhancing the reusability of  the materials. The following sections describe the oxidation mechanism  for pure TaC and HfC and their solid solutions.  3.2.1. Oxidation behaviors of pure TaC and HfC Fig. 3a-c shows the oxide scales of pure HfC after 60 s, 180 s, and  300 s exposure. The cross-section of HfC oxidized samples was porous,  and only a single layer oxide was observed. Samples  from different  exposure duration did not show obvious mechanistic differences except  for  additional oxide  layer  thickness. The oxide  layers  for  the  three  samples were uniform. Oxidized grains  could be  seen in the  cross section, suggesting that the oxidation initiated at the grain boundaries.  The cracks were inter-connected, which provided a pathway for  the  gaseous products without disrupting the oxide layers. In the conven tional  isothermal oxidation studies on the oxidation behavior of HfC,  one of the important reasons for HfC has superior oxidation resistance  is  the  formation of  a dense,  crack-free  oxycarbide  layer, which is  considered as a protective layer. However,  in the present study, such  oxycarbide layer is barely visible; only few micron thick, the dense layer  can be seen, as  shown in Fig. 3b. HfC can dissolve oxygen without  turning into HfO2, so the oxidation process of HfC always begins with the adsorption and diffusion of oxygen into the lattice. However, in the  present  study,  high  temperature  and  high  velocity  accelerate  the  diffusion process. Hence,  the oxycarbide phase became unstable and  transformed into HfO2. The cross-section images of oxidized TaC shown in Fig. 3d-f. The  oxide  scales were  smoother  compared to the oxidized HfC samples  suggesting severe melting during the oxidation tests. Giant cracks (up  to 10 µm opening) were  found in the oxide  scales of oxidized TaC  samples due to the gaseous products during the oxidation and thermal  stress  upon  cooling.  Spallation was  noticed  in  all  TaC  samples,  especially in the 300 s test.  The XRD analysis was performed on the  front  surfaces  of  the  Table 1  Densification and average grain sizes of HfC, TaC, and HfC/TaC solid solutions.  Name  Pellet density (g/cm3)  Densification (%)  Avg. grain size (µm)  Microhardness (GPa)  Elastic modulus (GPa)  Indentation toughness (MPa m1/2)  Pure TaC  14.14  96.7  6.8 ± 1.4  12.27 ± 0.87  331.67 ± 4.23  4.56 ± 0.52  T80H20  13.85  97.8  6.2 ± 2.1  16.39 ± 0.86  443.24 ± 23.65  4.58 ± 1.06  T50H50  13.26  98.2  3.8 ± 1.2  17.15 ± 1.1  523.82 ± 7.03  6.03 ± 0.70  T20H80  12.68  98.8  3.1 ± 1.1  19.06 ± 0.27  577.30 ± 6.04  5.51 ± 0.56  Pure HfC  12.21  98.5  2.3 ± 0.7  18.46 ± 0.22  360.86 ± 29.53  3.39 ± 0.97  Fig. 2. Oxide layer thickness comparison between five samples after exposure to high temperature plasma for 60 s, 180 s, and 300 s, respectively.  Table 2  Carbide recession rate of TaC, HfC and TaC/HfC solid solutions.  Sample  TaC  T80H20  T50H50  T20H80  HfC  Carbide recession rate (μm/s)  1.51  0.82  0.26  0.91  1.19  C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  14800  \\x0c\", \"samples  for  all  three  testing durations. The  oxidized HfC samples  showed a single phase: HfO2 Fig. 4a. The XRD patterns for TaC oxidized samples are  for all  three-time durations  shown in  shown in  Fig. 4b and C. The 60 s and 180 s samples are identical and indexed as  orthorhombic Ta2O5. However, addition to the orthorhombic Ta2O5 phase, triclinic Ta2O5 was also detected. The triclinic Ta2O5 is reportedly a high-temperature form of tantalum pentoxide [12,19]. The transformation of tantalum oxide  in the 300 s test sample (Fig. 4c),  in  from high-temperature form to low-temperature form was sluggish, so  the high-temperature form of the oxide can be retained when samples  were  quenched. Delamination and spallation were  severe  in 300 s  oxidized sample, which provided more pathway to air and resulted in  high cooling rate.  3.2.2. Oxidation mechanisms of TaC-HfC solid solutions  Fig. 5 displays the cross-section images of  the oxidized HfC/TaC  solid solution samples. The oxide scales remained intact without any (Fig. 5a-c) of  spallation or delamination. The  cross-section images  T20H80 samples showed a single oxide layer with a uniform thickness.  Although cracks covered the oxide scales, they were denser than those  for pure HfC. The thickness of  the oxide scales  in T20H80 samples  reduced by more than 50% as compared to pure HfC oxidation. From  the cross-section images,  the grain shapes could still be seen, so the  grain boundaries were the preferential diffusion pathway as well  for  T20H80 samples. Since the testing temperature was high enough, the  resultant HfO2 sintering aid. As a result,  could  be  sintered with  liquid  Ta2O5 in T20H80 samples were  serving  as  a  the oxide scales  denser than those in pure HfC samples.  The cross-section pictures of the T50H50 post-oxidation samples shown in Fig. 5d-f presented a very dense oxide scale. The oxide scales  and residual carbides had distinct interfaces, and no other layers could  be found in the scales. Some cracks still can be spotted inside the oxide  scales,  but  crack-healing  can  also  be  observed  as  the  oxidation  proceeded. Such thin oxide  scales  indicating  a protective oxidation  mechanism was dominant during the T50H50 oxidation.  Although the oxide scale in T80H20 samples were full of cracks, no sign of delamination was observed (Fig. 5g-i).  In the 300 s  testing  sample  (Fig. 5i),  signs of  liquid infiltration at  few locations  at  the  interface are spotted (marked with arrows)  indicating massive liquid  phase generated during the testing. Such signs of liquid attacking could  also be found in T50H50 samples, but T80H20 showed much severe  attacking, especially in 300 s  testing. This might be due to the low  melting point of Ta2O5 from the increased amount of TaC. The relative content of Ta2O5 increased from 26% in T50H50 sample to 50% in T80H20 sample  after  300 s  oxidation  tests,  as  obtained  by  the  integrated intensity method from the x-day diffraction data. Another  interesting  characteristic  of T80H20  cross-section  images was  the  significant  crack-healing. This phenomenon could also be explained  by the liquid phase formation, which will be elaborated in the following  text.  The XRD patterns for T20H80 samples after oxidation are shown in  Fig.  6a. With  20 vol% of TaC addition,  the  post-oxidation phases  showed a mixture of HfO2 and Hf6Ta2O17. The phase's matched the early oxidation work on the oxidation of HfC-TaC solid  information  solutions [16,17,19,20]. There is no substantial difference between the  samples for the three-time durations. The XRD patterns for T50H50  samples were the same for the three-time durations and displayed in  Fig. 6b. The overall patterns were similar to the ones for T20H80 samples. The major difference was at 22-23° where Ta2O5 peak started to appear. This is due to the increased amount of TaC addition  Fig. 3. The cross-sectional  images of the oxide scales of the pure carbides (a-c) pure HfC: 60 s, 180 s, and 300 s; (d-f) pure TaC: 60 s, 180 s, and 300 s.  C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  14801  \\x0c\", 'compared to the T20H80 samples. The formation of Hf6Ta2O17 is the result of the reaction between Ta2O5 and HfO2 as shown in the  Reaction (1).  Ta O  +6 HfO  Hf Ta O  →  2  5  2  6  2  17  (1)  At  the extreme testing conditions,  individual Ta2O5 and HfO2 can  still be detected, suggesting the reaction was blocked by separating the  Ta2O5 and HfO2. oxidation are shown in Fig. 6c. At first glance,  The  XRD patterns  for  T80H20  samples  after  the XRD peaks for the  three-time durations were  identical. All  the peaks were  indexed as  orthorhombic Ta2O5, the low-temperature form of tantalum pentoxide. No presence of HfO2 was detected due to the relatively low amount of HfO2 Formed. A closer look at 180 s and 300 s tests XRD pattern, (shown as marked by arrows in Fig. 6c),  tiny peaks started to appear  and are indexed as Hf6Ta2O17. The formation of Hf6Ta2O17 requires more HfO2 than Ta2O5 according to reaction (1). In the 1 min test, there was not enough HfO2 formation, so no Hf6Ta2O17 was detected. Only after more than 180 s exposure when enough HfO2 had formed, Hf6Ta2O17 started appearing. The oxidation of pure HfC and pure TaC have been investigated by  different researchers, but largely restricted to isothermal oxidation conditions inside a conventional furnace [21-24]. Due to the different  testing conditions during isothermal oxidation (with no strong air flow  and relative low testing temperature that  is no high enough to melt  Ta2O5), applied to the current study. Fig. 7 presents the oxidation mechanism  the  proposed mechanisms  and models  cannot  be  directly  for the oxidation of pure HfC samples based on the test conditions in  the  present  study.  The HfC  samples were  sintered  pellets  from  powders. The grain boundaries are  the preferred diffusion pathway  due to their high defects concentration. The oxidation process begins  with the absorption of oxygen,  followed by the diffusion of oxygen  through the grain boundaries and formed HfO2 and grain surfaces (Step 1). When the HfC pellets are exposed to the high-temperature  oxygen flow,  the initial oxygen absorption and oxidation are acceler ated  due  to  the  high  temperature  and  velocity  as  compared  to  isothermal oxidation test  in a conventional  furnace. The HfC grains  at the surface rapidly oxidized and formed HfO2. The grain boundaries became vulnerable due to the high defect concentration and the high  vapor pressure from the gaseous products. As a result, cracks occur and  create new pathways  for  oxygen diffusing deeper  into un-oxidized  carbides, starting a new cycle of oxidation (Step 2). With the near sonic  speed, the oxygen was pushed through those formed cracks instead of  diffuse through the grains to oxidize the underlying carbide. As shown in Fig. 3a-c,  the oxide scale in pure HfC samples are granular; each  grain is isolated by intermediate gaps. The grain size of the oxide scale  was found to vary due to the localized sintering. Only a very thin dense  layer  (< 5 µm) was  seen which matched the literature description of  oxycarbide phase suggesting the high oxygen potential throughout the  oxide scale [5,25]. The sintering of HfO2 densified the oxide scales, and the proper “Pilling-Bedworth ratio” (Voxide/Vcarbide = 1.39) allows good adherence to the carbide (Step 3). The sintering temperature for a  material is considered the two third of its melting point and the melting  point  of HfO2 temperature was close to its melting point, which could provide enough  is  around  2800 °C.  The  present  oxidation  testing  driving  force  for  the  sintering  of HfO2. As result, HfO2 obvious spallation was noticed under  a  scales  remained intact,  and no  the  extreme conditions.  Pure tantalum carbide exhibited poor oxidation resistance in the  present study. Not only it had the thickest oxide layer, but  its oxide  scales were also peeled off from the carbide surfaces. Similar  to the  oxidation  of  hafnium carbide,  the  oxidation  process  in  tantalum  carbide began at  the grain boundaries. The oxygen diffusion/penetra tion process was akin to the process in the HfC; cracks and gaps were  formed at the grain boundaries and allowed oxidation into the deeper  carbide. The difference was  the  oxygen diffusion pathway  into the  grains.  In the  case of TaC,  the  formed Ta2O5 on the grain surface rapidly melted under the high-temperature environment and covered  the TaC grains. The liquid phase formation hindered the oxygen lattice  diffusion, but  the liquid layer was  far  from being protective.  It was  pointed out  that Ta2O5 would react with TaC and further oxidize the underneath carbides [26]. The formed Ta2O5 cannot withstand the environment and  high  temperature  (2800 °C)  liquefies  during  the  Fig. 4. XRD patterns of (a) HfC and TaC oxidized samples (b) 60 and 180 s, (c) 300 s.  C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  14802  \\x0c', 'C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  Fig. 5. The cross-sectional images of the oxide scales of the HfC/TaC solid solutions, (a-c) T20H80: 60 s, 180 s, and 300 s; (d-f) T50H50: 60 s, 180 s, and 300 s; (g-i) T80H20: 60 s,  180 s, and 300 s (arrows showing liquid attacking).  oxidation process. The gaseous products generated can easily disturb  and lift  the liquid phase oxide scales, causing massive spallation and  cracking. Although liquid phase is considered beneficial  to the oxida tion resistance, the liquid phase itself should not react with the parent  material. Also without the solid phase structure, the liquid phase would  be prone to ablation under the high gas blow.  TaC-HfC solid solutions had much better oxidation resistance than  their monolithic carbides. In the case of T50H50, the oxide scales were  not only thin but also highly dense. The thickness of T50H50 oxide  that  the Ta2O5 has been consumed by the reaction. The formation of one mole of Hf6Ta2O17 requires one mole of Ta2O5 and six moles of HfO2 so that no Ta2O5 would remain with the presence of HfO2. Additionally, multiple reports suggest that the formation of Hf6Ta2O17 requires at least 10 h and the temperature exceeding 1200 °C [27,28].  Although  the  current  testing  conditions  accelerated  the  reaction,  can  only  Hf6Ta2O17 of Ta2O5 HfO2. The Hf6Ta2O17 formed was a protective layer against oxidation, so the thickness of the oxide scale had reduced up to 65%. However, the  formation  formed  after  and  the  be  scale was 1/10 of the thickness of pure TaC 300 s oxidation, 1/6 of the  amount of  liquid phase formed was not enough to form a continuous  thickness of pure HfC 300 s oxidation. The oxidation mechanisms for  TaC-HfC solid solutions could be deduced from the oxidation mechan liquid layer and protect the T20H80 samples cross-section images (Fig. 5a-c) showed that  from oxidation. The  though the oxide scales  isms for oxidation of monoclinic TaC and HfC. As shown in Fig. 8, the  were not as granular as the ones in pure HfC samples, the large cracks  first step is the same as pure HfC oxidation case. Liquid phase starts to  and big gaps still could be seen. Some minor cracks along the grain  appear from the molten Ta2O5 (Step 2). The oxidation process shifts from oxygen attack in the pure HfC cases to liquid phase attack to the  boundaries indicate that  the same oxidation mechanism is prominent  in both  pure HfC and T20H80  samples. Additionally,  oxide  scale  underlying carbides (Step 3). The HfO2 would serve as a solid scaffold and anchor the liquid phase that can lead to cracking healing (Steps 4  sintering occurred, especially in the presence of liquid phase, explained  some large grains found in T20H80 oxide scales.  and 5).  When 50 vol% of the TaC was added to T50H50 samples, the more  In the T20H80 sample,  the 20-vol% TaC was added. During the  liquid phase was generated during the oxidation. The beginning of the  oxidation, HfO2 and Ta2O5 were formed at the surfaces of the top layer grains. Due to the extreme testing conditions, especially the high  temperature,  the formed Ta2O5 melted and covered on HfO2 surface. like in the pure TaC oxidation case, the formed liquid Ta2O5 was not a protective layer against oxidation. It could react with both  However,  analysis  TaC and HfC and form tantalum suboxide  and HfO2. samples, only HfO2 Hf6Ta2O17 were detected without the presence of Ta2O5, which implies  post-oxidation  In the XRD  T20H80  and  the  on  oxidation was the same as the pure HfC and the T20H80 samples. The  XRD patterns showed the mixture of Ta2O5, HfO2, and Hf6Ta2O17. As discussed in the earlier, the Ta2O5 and HfO2 could not co-exist, the Ta2O5 should be consumed into the formation of Hf6Ta2O17 at such high temperature. However, the XRD resulted suggest otherwise. The  only explanation was the formed Ta2O5 and HfO2 were separated. In the T50H50 samples, the Ta2O5 formed on the HfO2 melted and formed Hf6Ta2O17 as shown schematically in Fig. 9.  14803  \\x0c', 'C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  Fig. 6. XRD patterns of post-oxidation samples. (a) T20H80, (b) T50H50, and (c) T80H20.  Fig. 7. Schematic of the oxidation mechanisms for pure HfC samples. (Step 1: oxygen attacks under harsh environment, Step 2: oxygen penetrates via grain boundaries left by gaseous  products from oxidation process, Step 3: porous oxide scale forms).  14804  \\x0c', 'C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  Fig. 8. Schematic of the oxidation mechanism for TaC-HfC solid solutions. (Step 1: oxygen attacks under harsh environment, Step 2: oxide scale formed melts, Step 3: oxidation process  through liquid phase attacking, Step 4: oxide scale formed with large cracks, Step 5: crack healing due to the molten oxides).  Fig. 9. Illustration of the formation of Hf6Ta2O17 that separates the molten Ta2O5 and  HfO2.  The formed Hf6Ta2O17 did not react with underlying materials and also separate the Ta2O5 and HfO2. The abundant liquid phase quickly eased the stress at the grain boundaries, so the oxide scale was much  denser  compared  to  the  other  samples. Moreover,  enough  liquid  formed during the oxidation which formed a continuous liquid layer.  Two different  liquid phases exist  in the T50H50 samples resulted in  high viscosity fluid, which could withstand the high gas flow. The  formed HfO2 also provided a scaffold structure that the oxide scales. Such scaffold structure was able to observe from the  further stabilized  FIB sectioned surface from the oxide scale shown in Fig. 10a. The milled surface displayed a “ridge” and “valley”  structure due to the  different milling resistances from Hf6Ta2O17, HfO2 and Ta2O5. Since the resultant Hf6Ta2O17 and HfO2 did not melt during the oxidation testing, they had well-defined crystal structures that withstood the milling process and resulted in “ridge” structure. However Ta2O5, on the other hand, melted during the oxidation process. It was more likely  to retain amorphous phase upon cooling that was easily milled in the FIB process. The EDS point analysis conducted on the “valley” (Ta2O5 rich) and “ridge” (HfO2 and Hf6Ta2O17 rich) is shown in Fig. 10b. “valley” Points 1 and 3 were performed on the areas while 2 was conducted on the “ridge” area. Since the FIB milling and EDS analysis  was conducted on T50H50 samples, both Ta and Hf should have the  same atomic percentage. At locations 1 and 3, the EDS results indicated  Ta-rich areas where point 2 showed a Hf-rich area. Based on these observations and milling features, it is safe to conclude that the “valley”  area is Ta2O5 dominated which is the liquid phase in the proposed solid  Fig. 10. (a) Scaffold structure in the solid solution oxide layer obtained from the FIB  sectioned solid solution oxide layer. (b) EDS analysis results on the “ridge” and “valley”  areas.  scaffold and liquid phase structure. In the “ridge” area, the Hf atomic  content  is not significantly greater than Ta. This can be explained by  of  The  the  solidification  Ta2O5. un-melted HfO2 and Hf6Ta2O17 provided the nucleation sites for the molten Ta2O5 and resulted into crystallized Ta2O5 around the “ridge” areas. The morphology of T80H20 60 s test sample (Fig. 5g) was similar to  the one of the pure TaC (Fig. 3d), but the thickness of the oxide scale of  T80H20 was only 45% of  the thickness of pure TaC oxide scale. The  14805  \\x0c', \"main difference was  the 20 vol% of HfC in the  system. Due  to the  relatively low amount of HfO2, HfO2 was not detected by XRD. The HfO2 secured the oxide scale from ablation, and the oxide scale remained intact. Fig. 11 shows that a single grain was attacked and  lifted by the liquid phase indicating the liquid phase prevented the  oxygen diffusion, and the oxidation of T80H20 was  from the liquid  attack to the carbides. In the 180 s and 300 s test samples (Fig. 5h and  i), crack-healing was again noticed, which also verified the existence of  HfO2 structure was not noticed in pure TaC samples,  that  hosted  and  stabilized  the  liquid  phase.  Such  scaffold  so no crack-healing  was noticed. In the 300 s test,  the Hf6Ta2O17 peaks started to appear due to the HfO2 formation of a large amount of HfO2 generated during the oxidation tests. The liquid layer became more protective against the  oxidation,  and further  crack-healing occurred as  seen in the  cross section images.  4. Summary  The oxidation studies in the present work were conducted on TaC,  HfC and their solid solutions (T80H20, T50H50, and T20H80) using  high-speed and high-temperature plasma  jet. The  results  exhibited  better oxidation resistance  in TaC-HfC solid solution samples  than  their constituent carbides. The T50H50 has the thinnest oxide layers  for three-time durations, its oxide scale thickness is 1/10 of TaC's, and  1/6 of HfC's  after 300 s  exposure. The main oxidation mechanism  shifted from the diffusion through the oxide scale to the attacking on  the grain boundaries due to the harsh testing conditions. The addition  of TaC resulted in a low melting point Ta2O5 that further changed the oxidation mechanism to liquid phase attacking the grain boundaries.  Hf6Ta2O17 is the reaction product between HfO2 and Ta2O5 was found in the oxide scales of the solid solution samples, which is the reason for  better  oxidation resistance. More  importantly,  a  solid scaffold and  liquid phase structure was achieved in the oxide scales of  the solid  solution samples, which not only provided improved oxidation resis tance but also enhanced the mechanical integrity of the oxide scales by  healing the cracks during the oxidation process. This work provides an  insight into the carbide based solid solutions as candidate materials for  use on the hypersonic vehicles that require surviving under the extreme  conditions.  Acknowledgement  Cheng Zhang would like to thank University Graduate School at  Florida  International  University  (FIU)  for  Dissertation  Year  Fellowship. Authors  also acknowledge  the  characterization facilities  at Advanced Materials Engineering Research Institute (AMERI) at FIU.  References  [1] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. Zhou, Ultra-High Temperature  Ceramics: Materials for Extreme Environment Applications, John Wiley & Sons,  Inc., Hoboken, New Jersey, 2014.  [2]  L.E. 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Zhang, Effect of SiC content on the ablation and  [17]  [18]  oxidation behavior of ZrB2-based ultra high temperature ceramic composites, Materials 6 (2013) 1730-1744. [19] A. Paul, J.G.P. Binner, B. Vaidhyanathan, A.C.J. Heaton, P.M. Brown, Oxyacetylene  torch testing and microstructural characterization of tantalum carbide, J. Microsc. 250 (2013) 122-129. Y. Wang, X. Xiong, G. Li, H. Liu, Z. Chen, W. Sun, et al., Preparation and ablation  [20]  properties of Hf(Ta)C co-deposition coating for carbon/carbon composites, Corros. Sci. 66 (2013) 177-182. [21] M. Desmaison-Brut, N. Alexandre, J. Desmaison, Comparison of the oxidation behavior of two dense hot isostatically pressed tantalum carbide (TaC and Ta2C) materials, J. Eur. Ceram. Soc. 17 (1997) 1325-1334. S. Shimada, K. Nakajima, M. Inagaki, Oxidation of single crystals of hafnium  [22]  carbide in a temperature range of 600 °C to 900 °C, J. Am. Ceram. Soc. 80 (1997) 1749-1756. [23] D. Liu, J. Deng, Y. Jin, C. He, Adsorption of atomic oxygen on HfC and TaC (110) surface from first principles, Appl. Surf. Sci. 261 (2012) 214-218. S. Shimade, Interfacial reaction on oxidation of carbides with formation of carbon, Solid State Ion. 141-142 (2001) 99-104. [25] R. Savino, M. De Stefano Fumo, L. Silvestroni, D. Sciti, Arc-jet testing on HfB2 and HfC-based ultra-high temperature ceramic materials, J. Eur. Ceram. Soc. 28 (2008) 1899-1907. [26] A. Lashtabeg, M. Smart, D. Riley, A. Gillen, J. Drennan, The effect of extreme  [24]  temperature in an oxidising atmosphere on dense tantalum carbide (TaC), J. Mater. Sci. 48 (2013) 258-264. [27] M. Li, Q. Xu, S. Zhu, L. Wang, F. Wang, Preparation and thermal expansion of Hf6Ta2O17 ceramic, Rare Met. Mater. Eng. 40 (2010) 612-614. [28] M. Li, Q. Xu, L. Wang, Preparation and thermal conductivity of Hf6Ta2O17 ceramic, Key Eng. Mater. 434-435 (2010) 459-461.  Fig. 11. The liquid phase is attacking a single grain through grain boundaries.  C. Zhang et al.  Ceramics International 43 (2017) 14798-14806  14806  \\x0c\"]"
},{
  "_id": 195,
  "PDF": "Oxidation, mechanical and thermal properties of hafnia–silicon carbide nanocomposites.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 34 (2014) 1783-1790  Oxidation, mechanical and thermal properties of hafnia-silicon carbide nanocomposites  Yutaka Shinoda a , David B. Marshall b , Rishi Raj c,∗  a Materials and Structures Laboratory, Tokyo Institute of Technology, 4529 Nagatsuta-cho, Midori-ku, Yokohama, Kanagawa 226-8503, Japan b Teledyne Scientific Company, 1049 Camino Dos Rios, Thousand Oaks, CA 91360, United States c Department of Mechanical Engineering, University of Colorado at Boulder, Boulder, CO 80309-0427, United States  Received 28 September 2013; received in revised form 8 December 2013; accepted 21 December 2013  Available online 18 January 2014  Abstract  Dense nanocomposites constituted  from 70/30 vol% of hafnia-silicon carbide and were prepared by  spark plasma  sintering. Silicon carbide suppresses grain growth. The  fracture strength of as prepared composites  is 400-600 MPa. Oxidation up  to 1600 C  in air  for 10 h has minor influence on the mechanical strength, which is ascribed to the dense nature of the oxidation scale. The high density of the oxidation scale is attributed to a volume increase when silicon carbide oxidizes and reacts with hafnia to form hafnium silicate. The composite has a thermal conductivity of −1 K −1 at room temperature. Design approaches for further enhancement of ultrahigh temperature properties of oxide/non-oxide composites 14 W m are discussed. © 2014 Elsevier Ltd. All rights reserved.     Keywords: Hafnia; Silicon carbide; Nanocomposites  1.   Introduction  While  there  is a  large body of  literature on ceramics made from composites of zirconium boride and silicon-carbide, there is much  less attention given  to dual phase composites made from oxides and silicon-carbide. The silicon carbide phase plays different  roles  in  these  two classes of materials.  In borides  it gives rise to silica during oxidation, providing passivation protection  to some extent.  In  the oxide-nonoxide composites  the silicon carbide phase is intended to impart thermal conductivity, and promote the evolution of hafnium silicate during oxidation, which, as shown  in  this article, yield a  robust oxidation scale with good mechanical properties. In oxidizing environments  the borides produce boron  trioxide (boria), while silicon carbide produces silica. Boria has low melting point (450-510 C), sublimes at 1500 −1 ). In comparison silica C, boils at 1860 C, and  is water soluble (22 g L melts at 1600-1725 C, does not sublime, has a boiling point of 2230 C, and has much lower solubility in water. Silica, having                 ∗  Corresponding author. Tel.: +1 303 492 1029.  E-mail address: rishi.raj@colorado.edu (R. Raj).  0955-2219/$ - see front matter © 2014 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.12.044  to oxygen,   is an effective passivation barrier  low permeability  to oxidation. Zirconium diboride when oxidized yields zirconia and  liquid boria, which flows and volatilizes leaving behind columnar channels that hasten oxidation.1 The addition of silicon-carbide oxidizes into silica which provides some degree of passivation. In  its presence a  silica  layer develops on  the  surface, underneath which a  layer of zirconia  immersed  in silica  is  formed. In between  this  layer and  the dense  substrate  is a  transition interlayer with partially depleted SiC.2 This multilayer oxide scale grows quickly,  reaching more  than 50  \\u242em at 1500 C  in 30 min,2 and several hundreds of micrometers near 1600 C  in 100 min.3 Above 1800 C and in low oxygen pressure environments, the silica passivation layer begins to volatilize into SiO.4 In aero-thermal, high velocity environments, the silica top layer is  lost completely  to volatilization and  is replaced by zirconia and a residue of carbon, with the oxidation structure lying underneath remaining  the same as  in static oxidation (with zirconia immersed in silica, etc.).5 In-situ X-ray diffraction studies of oxidation at  temperatures up  to 1650 C reveal crystalline phases of zirconia.6 Further additions of rare-earth oxides to the ZrB2-SiC composites produce zirconates  that alter  the phase constitution of                   \\x0c', '1784   Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790  Fig. 1. Scanning electron micrographs of the powders used in this study.  the oxide scale but do not appear to have a remarkable influence on the thickness of the oxide scale.7 The addition of zirconium silicide, which also promotes the formation of silica, appears to improve the oxidation resistance.8 Clearly the oxidation process of boride-carbide composites is complex. The HfO2 -SiC composites explored in this article are shown to have a straightforward pathway  for  the development of  the oxidation scale. The reaction with atmospheric oxygen produces hafnium silicate. SiC,  the minority phase, oxidizes  to produce silica which  then reacts with HfO2 to produce HfSiO4 . In  this way  the oxide  scale  evolves  into  a  composite of HfO2 and HfSiO4 . The escape of carbon-monoxide, a product of oxidation, leaves behind porosity, which surprisingly, is quite sparse and does not appear  to have a significant detrimental effect on the mechanical strength of  the specimens after oxidation. The addition of SiC imparts thermal conductivity to the composite. The formation of HfSiO4 in the HfO2-SiC composites stands in a remarkable contrast to the oxidation of the ZrB2-SiC which produce ZrO2 but, apparently, not ZrSiO4 . On  the other hand oxidation of HfB2-SiC composites have shown to form HfSiO4 composites.9 In one case10 only HfO2 has been  reported, but X-ray examination of  the oxidation scale was not  included  in this study. In  the  present  experiments  the  evolution  of HfSiO4 fortunate  since  the  reaction HfO2 (s) + SiC(s) + 1.5O2 (g)  HfSiO4 (s) + CO(g) produces a volume expansion of 13% among  →  is  the solid phases. The expansion  is presumably accommodated by  in-situ creep of  the oxide, which preserves  the density of the scale while  relieving  the compressive stress  from volume expansion.  2. Materials and processing     The  composites were  prepared  from  powders  of  hafnia, obtained  from American Elements Company,  and  ␤-silicon carbide from MTI Corporation. SEM micrographs from the powders are shown  in Fig. 1. In both  instances  the particle size  is less than 100 nm. The hafnia and silicon carbide powders were planetary ball milled in pure ethanol using SiC pot and balls at 300 rpm for 1 h and dried in air at 80 C. The powders were consolidated by Spark Plasma Sintering at temperatures ranging from 1450 C  to 2000 C at 50-100 MPa for 3-5 min  in nitrogen atmosphere. The microstructure of  the composite prepared at 2000 C  (50 MPa, 5 min)  is compared to  the pure hafnia polycrystal fabricated at 1800 C (100 MPa, 3 min) in Fig. 2. Most noteworthy is the large grain growth in the hafnia sample, which prevented  the sample from reaching full density. The SiC particles in the composite suppress grain growth 500 nm for hafnia and  200 nm for SiC. with a grain size of  The composite sample is fully dense. These values for the grain size were confirmed, at least approximately, from micrographs of the fracture surface of the composite sample.              Fig. 2. Microstructures of HfO2 -SiC composites, and HfO2 made by spark plasma sintering. Note the suppression of grain growth in the composite sample.  \\x0c', 'Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790   1785  Fig. 3. Left: The thickness of the oxide scale as a function of the temperature of oxidation. The samples were oxidized in air for 10 h. Right: A comparison of present  with data for zirconium-diboride based composites from the literature. The thickness of the oxide scale as a function of the temperature of oxidation. The data for  the borides have been extrapolated to 10 h using power law equation with an exponent of 2 (parabolic oxidation) or 3. The error bars for the borides represent these  two exponents.  3. Properties        ×  ×  Samples prepared by SPS were cut, ground and mirror polished with 1  \\u242em diamond paste,  into  the shape of  rectangular bars of dimensions 1 mm   1.3 mm   12 mm for 4-point bend testing. Another set of specimens, also mirror polished, were oxidized  in ambient air for 10 h at  temperatures ranging from 1200 C  to 1600 C. These samples were  tested  in  four-point bending with  the as-grown oxidation scale  in place. The oxide scale on  these  samples was characterized by cross-sectional SEM and X-ray diffraction. The thermal diffusivity of the composite was determined by the  laser-flash  technique, using a commercial  test  instrument (Netzsch, LFA 427). The test samples were in the form of disks of diameter 10 mm and thickness 1.17 mm, with surfaces polished as described above then coated with a thin layer of carbon. Measurements were made at selected  temperatures between  room temperature and 1400 C. The sample was heated  in argon at a  rate of 20 /min between  sequential measurement  temperatures. At each measurement point  the  temperature was held constant for a period of 10 min during which  three diffusivity measurements were taken. In  the following section  the oxidation behavior  is described first, which is followed by the measurements of the mechanical strength and thermal conductivity.        3.1. Oxidation behavior        Samples  were  oxidized  for  10 h  in  ambient  air  at 1200-1600 C  in steps of 100 C. The  thickness of  the oxide scale for each specimen was measured. X-ray diffraction spectra from the surface of the oxidized specimen were obtained. All samples gained weight by an amount  that ranged from 0.03% −2 ) at 1200 −2 ) at (or 0.59 mg cm C,  to 2.04%  (or 3.98 mg cm 1600 C. The  thicknesses of  the oxide scale, which were measured directly from cross-sections, are described below. The scale  thickness  from present experiments  for samples oxidized  in air  for 10 h are given on  the  left  in Fig. 3. They are compared to data for zirconium diboride based materials on        t   n   2,  =  hn =  the  right  (the diboride data  in Fig. 3 were extrapolated  from literature values to 10 h using the equation  kt , where h is the  thickness of  the oxide scale, k  is  the rate constant, and  is the  time. Two values of   3 were used, which resulted  in the error bars shown  in Fig. 3 on  the right). The variability  in the data for ZrB2 materials is notable.2,3,6,7,11-15 In general the HfO2-SiC composites show better resistance to oxidation than the borides. For example, the hafnia composites develop a 127  thick scale  in 10 h at 1600 C, whereas  in the case of borides the scale has been shown to grow to several hundred micrometers at 1600 C  in  just 100 min.2,3 While  the data for the present HfO2-SiC composites need to be reproduced in other laboratories, they appear to bear consistency, which can enhance  the predictability of  the performance of  this class of ultrahigh temperature ceramics.  \\u242em         3.2. Microstructure of the oxide scale     The microstructure  of  the  oxidation  scale  obtained  after exposing the HfO2-SiC composite to 1600 C for 10 h in ambient air  is  shown  in Fig. 4. The microstructure changes with the distance from  the surface. The compositions of  the various phases were identified by EDS. In regions close to the interface, the EDS spectra indicate the presence of SiC and SiO2 (as well as hafnia). With increasing distance from the interface, the fraction of remaining SiC decreases to zero and the silica reacts with hafnia to form hafnon near the surface. The compositions close to the interface represent the early stage of oxidation, while the near-surface regions represent the later stages. The oxidation reaction requires  the escape of CO(g) which accounts for  the presence of pores within  the oxide scale. The upper bound pressure of CO within these pores is calculated in Appendix. These estimates are too high to be contained within closed pores.  It  is  therefore postulated  that  the pores provide channels for CO to escape into the atmosphere. The volume change from the reactions leading up to the evolution of hafnium silicate can be estimated at least approximately. The net reaction is the conversion of SiC and hafnia into hafnium silicate. The densities of  these constituents are 3.21, 9.68 and      \\x0c', '1786   Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790  Fig. 4. EDS spectra showing the composition of the phases observed at various locations in the oxide scale.  −3 respectively (the density for the silicate was calcu6.97 g cm lated from its lattice parameter), while their molecular weights −1 . These values show  are 40, 210, and 271 g mol that HfSiO4 has 13.8% greater volume  that  the sum of  the SiC and HfO2 from which  it  is produced. Thus,  it  is  inferred  that  the oxide scale grows under a compressive stress, which  is relaxed continuously by creep as the scale grows. The final outcome of this pathway  is  to produce a dense oxide scale, except for  the CO pores, as shown in Fig. 5. X-ray diffraction  spectra were obtained  from  the  surfaces of  the oxidized  specimens  to determine  the phases present. Results from samples that had been oxidized for 10 h are given in Fig. 5. There  is a difference between specimens oxidized at 1500-1600 C  than  those oxidized at  lower  temperatures. The     specimens oxidized at  the higher  temperatures show  the presence of hafnium silicate and hafnia (monoclinic), while the data from  lower  temperature oxidation show mostly hafnia (monoclinic) and silica  (cristobalite)  (the peaks at about 29, 39, 43, 48 and 49 , that are especially visible for specimens from lower temperatures, are from the clay used to fix the specimens).     3.3. Mechanical properties  The fracture strengths of the as-sintered and oxidized materials were measured in 4-point bending using a universal testing machine (AG-I, Shimadzu Co., Kyoto, Japan) operating at constant crosshead speed of 0.5 mm/min  (stress  rate 180 MPa/s).  \\x0c', 'Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790   1787  Fig. 5. X-ray diffraction of the oxidation scale formed at different temperatures. Note the presence of HfSiO4 , and the absence of cristobalite at the highest temperature.  All measurements were made at room temperature. The bending strength was calculated from  −  l)  = 3P (L  σ  2wt 2  l   where P  is  the maximum  load, L  (=10 mm)  (=3.3 mm) are the distances between  the outer and  inner  loading points, w (=1.3 mm) is the specimen width, and t (=1 mm) is the specimen thickness. The results are given  in Fig. 6. The error bars represent  the spread  in data obtained  from  three specimens  tested  for each oxidation condition. The average strength of the as consolidated  Fig. 6. Fracture strength of as-sintered and oxidized samples (in air at 10 h at  various temperatures).  composites was 660 MPa. The  strength of  the oxidized  samples varied  from 410 MPa  to 570 MPa, with  the exception of sample oxidized at 1200 C which gave a higher value than the as-sintered specimens. The reason for this higher strength is not understood. The retention of mechanical strength  is attributed to the dense structure of the oxide layer.     3.4. Thermal conductivity  −1  The measured values of thermal diffusivity of the composite are shown  in Table 1. Each value  is  the average of  three measurements, which varied by no more  than 0.02 mm2 s from the mean. The  thermal conductivity was calculated  from  the diffusivity data using values of specific heat and density calculated for  the composite from  literature values for SiC16 and HfO2 ,17 on the basis of rule-of-mixtures. The values of specific heat and thermal conductivity are shown in Table 1. The thermal conductivity  is plotted as a  function of  temperature  in Fig. 7. The  conductivity of  the  composite was  −1 K −1 ) at both room  substantially higher than  that of HfO2 −1 K −1 ) and at high temperature (5.6 W m (1.2 W m temperature −1 K −1 ). (13.1 W m The thermal conductivity of composites where high conductivity particles are dispersed  in an  insulating matrix depends on three parameters: the volume fraction, the thermal boundary resistance, also known as the Kapitza resistance of the interface between these phases, and the interconnectivity of the conducting phase. Rigorous solutions for the case where the particles are not interconnected are available.18 Here we give the result when the Kapitza resistance is negligible (meaning   0 in Ref. 18), and when the thermal conductivity of the particle is much greater  ≈  α          \\x0c', '1788   Table 1  Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790  Thermal conductivity of the composite.  T   C  Diffusivity,  mm2 s  −1  α  Cp (HfO2 )a −1 K J kg  −1  Hafnia-30 vol% SiC composite  Cp (SiC)b −1 K J kg  −1  Cp (composite) −1 K J kg  −1  Thermal conductivityc −1 K W m  −1  24   299   505.9   807.2   1103.1   1402   4.35   2.70   2.19   1.76   1.48   1.28   a Data from Ref. 16. b Data from Ref. 17. c Calculated from k =   285.6   348.9   360.9   369.0   373.6   377.2   719.6   964.5   1085.5   1188.6   1260.0   1318.0   339.9   425.9   451.5   471.4   484.4   494.8   13.1  10.2  8.7  7.3  6.4  5.6  α  ρ   Cp , where   ρ   is the composite density (8.87 g/cc).  immiscible,  the following observations suggest  that  they form strong,  chemically bonded  interfaces between  them:  (i) The presence of SiC as a second phase  retards grain growth during densification -  this would be possible only  if diffusion along the interfaces was slow which indicates interfaces of low energy.19 (ii) The fracture strength measurement suggests a good mechanical strength of  the oxide-nonoxide  interface. (iii) The thermal conductivity measurements suggest  low  resistance  to phonon transmission across the interfaces, which further imply that interfacial bonds possessed high stiffness. The hafnia-silicon carbide nanocomposites apparently have some  surprising properties.  In addition  to high  strength and high  thermal conductivity  the evolution of  the hafnium  silicate appears to impart mechanical strength to oxidized samples. The volume expansion as hafnia and silicon carbide react with oxygen  to create hafnium  silicate could have contributed  to the  integrity of  the oxidation  scale as well. Whether or not the evolution of hafnium silicate at  the surface of  the oxidation  scale provides additional  resistance  to  further oxidation remains unclear. This point  is being currently  investigated  in our laboratory. The present work raises  interesting new questions. (i) What is the atomic structure and properties of oxide-nonoxide interfaces? The  present work  suggests  these  interfaces  to  have unusual strength and stiffness, and low thermal boundary resistance. (ii) How can the grain size and the volume fractions of the constituents be adjusted to maximize the thermal conductivity? Measurement of the thermal conductivity of 30 vol% SiC composites gives values that are higher than predicted from models for mixtures.  It  is possible  that higher concentrations of SiC would produce  further  increase  in  thermal conductivity.  (iii) What can be predicted about the evolution of the oxidation scale? This is an interesting question. Here, we find that hafnium silicate is the end product, which along with residual hafnia forms the  two-phase structure of  the oxide scale. The constitution of the oxide scale can be adjusted  to different fractions of hafnon and hafnia. How  this microstructure affects  further oxidation, creep resistance and mechanical strength remains an interesting question. Finally we note that lifetime under tensile loading in oxidizing environments needs  to be  investigated, where creep crack growth  is  likely  to be  the failure mechanism. The deformation  Fig. 7. The thermal conductivity of the composite.  than that of the matrix. Two cases, one for low volume fraction and the other for high volume fraction have been solved.18 Here we consider the high volume fraction case which gives the ratio of the thermal conductivity of the composite,  κC , and that of the matrix,  κm , to be given by:  =  κC  κm  1  −  (1   f   )3  (1)  f   the  of    0.3   =  f  −1 ,11  κC =  κm =  phase. equal  where  is  the  volume  fraction  of  conducting  Assuming  the  thermal  conductivity  hafnia  to  be  −1 K to   1.2 W m and  gives  that  −1 K −1 . This value  2.9 W m is  far  lower  than  the experimen−1 K −1 . This difference  tal value of 13.1 W m implies  that  the silicon carbide particles had significant interconnectivity in the composite. More  importantly,  it also  implies  that  the  thermal boundary resistance of the HfO2 -SiC interface must have been low. A  low  interfacial  thermal  resistance  implies high elastic stiffness of the bonds across the interface, which then translates into strong interfacial strength, and thence to high resistance to fracture.  4. Discussion  The measurements of the oxidation behavior, the mechanical strength and  thermal conductivity of HfO2-30 vol% SiC composites show promise for oxide-nonoxide ceramics as ultrahigh temperature materials. Although  the  two phases are mutually        \\x0c', 'Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790   1789  and stress relaxation of  the dual phase hafnium silicate-hafnia microstructure near  the crack  tip would be an  interesting point to keep in mind.  5. Summary  Two phase SiC-HfO2 nanocomposites are discovered to have the following unusual properties:  (i) The presence of SiC particles  restricts grain growth during  hot  consolidation,  leading  to  nanoscale  two  phase microstructure which cannot be obtained  in  single phase hafnia materials. (ii) The  nancomposites  have  strengths  in  the  400-650 MPa range. The strength remains essentially constant after oxidation for 10 h at 1600 C in ambient air. (iii) The oxidation of  the nanocomposites  leads  to  the evolution of a dense  top scale constituted from hafnium-silicate and hafnia. The scale remains dense after oxidation which is attributed  to volume expansion as SiC converts  to SiO2 which then reacts with HfO2 to form hafnium silicate. The compressive stress developed by  the oxidation  reaction  is expected to have been continuously relieved by in-situ creep. (iv) The oxide scale remains dense and crack free when cooled down to room temperature which spells an absence of tensile stresses that could have been generated by a mismatch in thermal expansion. The possibility of compressive residual  stresses, however,  cannot be  ruled out. The  thermal expansion of  the oxide scale  is not known at  the present time. (v) There  is  good  evidence  that while  silicon-carbide  and hafnia  are  mutually  immiscible  in  bulk  form  they form  strong  interfaces  between  them.  This  point  can serve  as  a  topic  for  further  fundamental,  atomistic investigations.     Acknowledgements  This work was  supported  by  the US AFOSR  (Dr. Ali Sayir) and NASA (Dr. Anthony Calomino) under  the National Hypersonic  Science  Center  for Materials  and  Structures (AFOSR Contract No.  FA9550-09-1-0477). We  are  grateful  to Kalvis Terauds graduate  student  at  the University of Colorado, and Yuse Minoguchi, a  student at Tokyo  Institute of Technology  for coordinating  the work between  their  two Institutions.  Appendix A. Appendix  A.1. Calculation of the pressure within the CO pores in the oxidation scale  The oxidation of SiC produces CO(g) which must diffuse through  the oxidation scale. The micrographs  in Fig. 4 show the pores  to be present  throughout  the oxidation scale, which apparently serve as the escape route for carbon monoxide.  Here, we seek to calculate the upper bound for the CO pressure within the pores assuming that diffusion is prevented, and all of  the CO produced by oxidation of SiC accumulates  into the pores.  In  the analysis we use  the  following notation  (the superscript * refers to quantities after oxidation)  fSiC ,  fHfO2  volume fraction of SiC and HfO2 in the initial composite.  VC ,  V ∗  C  volume of the solid phases in the  composite before and after oxidation. −1 ) molecular weights (g mol  M SiC W  ,  M HfO2 W  ,  M HfSiO4 W  ρSiC ,  ρHfO2  ,  ρHfSiO4  densities in units of g cm  −3  nSiC ,  nHfO2  ,  n∗  SiO2  ,  n∗  HfO2  ,  n∗  HfSiO4  moles of various species in the  composite before and after oxidation.  The molar volumes of HfO2 and HfSiO4 are then given by  ¯vHfO2  = M HfO2 W ρHfO2  , ¯vHfSiO4  = M HfSiO4 W ρHfO2  ,  in units of cm3  (A1)  The pressure within  the pores  is obtained by converting  the quantities into moles, enforcing the reaction, and then converting the moles into volume using the values for the molar volume of the species. The reactions under consideration are  SiO2 +  SiC   +   1.5O2 (g)   =   CO(g)  and HfO2 +   SiO2 =   HfSiO4  (A2)  It therefore follows that the moles of the products are related to the moles of the reactants by:  nSiC =  n∗  SiO2  =  n∗  CO  n∗  SiO2  =  n∗  HfSiO4  n∗  HfO2  =  nHfO2  −  n∗  HfSiO4  (A3)  For  example,  in  the  last  equation,  the moles  of HfO2 remaining after oxidation are the difference between the original moles and the moles of HfSiO4 formed during oxidation. The volume of CO formed after oxidation at NTP is given by  V ∗  CO =  n∗  CO ×   2.24   ×   104  cm3 ,  (A4)  while the solid phases after oxidation will have the volume:  V ∗  C = ¯vHfSiO4  n∗  HfSiO4  + ¯vHfO2  n∗  HfO2  .  (A5)  Combining Eqs. (A1) and (A3)-(A5) leads to the following expression:  V ∗ V ∗  CO  C  = 2.24   ×   104  χ  ,  (A6)  where  χ  =   ( ¯vHfSiO4  − ¯vHfO2 )   + fHfO2 fSiC  · M SiC W M HfO2 W  · ρHfO2 ρSiC  · ¯vHfO2                                            \\x0c', '1790   Y. Shinoda et al. / Journal of the European Ceramic Society 34 (2014) 1783-1790  Note that the molar volumes as written in units of cm3 . We  convert  the CO gas pressure  at  ambient  temperature (1 atm)  to  the oxidation  temperature, T, using  the gas  law, giving:  V T  CO  V ∗  CO  =  T  300  ·  pT  CO  (A7)  where  CO is the pressure at the oxidation temperature, T, and CO is the actual volume of the CO pores at the oxidation temperature. Writing the total volume of the composite after oxidation, Ox as:  pT  V T  V ∗  V T  C =  V ∗  C +  V T  CO ,  and  the volume fraction of pores at  the oxidation  temperature as,  CO , the following final result for the CO pressure within the pores at the oxidation temperature is obtained  f T  pT  CO = T   ×   2.24   ×   104  300  · 1   −  f T  CO  f T  CO  · 1  χ  (A8)  All parameters on  the right hand side of  the above equation are known:   depends on the physical constants and the volume fraction of SiC  in  the original composite, T  is  the  temperature of oxidation, and  CO is the volume fraction of the pores in the composite after oxidation, which can be microstructurally determined. Recall that  CO is the upperbound value of the pressure since  the analysis assumed  that CO cannot diffuse out  into  the atmosphere. Substituting  oxidation   χ  f T  pT  the   following  at  1600  values  C   into   Eq.   (A8)  for      fHfO2  =   0.7,  fSiC =  0.3,  40 g mol 9.68 g cm  3.21 g cm 38.9 cm3 ,  21.7 cm3 gives  very  high  values for  the CO pressure at  the oxidation  temperature, which range from 67 kbar  to 32 kbar  for   0.05 to 0.10.  The  smaller value of  CO corresponds to the higher pressure. It is unlikely that the pores acquire such high pressure within them. A reasonable inference is the pores are interconnected to the surface  thereby serving as channels for outward migration of CO(g) into the atmosphere.  M SiC W  =  −1 ,  M HfO2 W −3 , ¯vHfSiO4  =   210 g mol  −1 ,  ρHfO2  =  −3 ,  ρSiC =  =   and ¯vHfO2  =  f T  CO =  f T  References  1. Parthasarathy TA, Rapp RA, Opeka M, Kerans RJ. Effects of phase change  and oxygen permeability   in oxide   scales on oxidation kinetics of ZrB2  and HfB2 . J Am Ceram Soc 2009;92(5):1079-86, http://dx.doi.org/10.1111/ j.1551-2916.2009.03031.x.  2. Rezaie A,   Fahrenholtz WG, Hilmas GE. Evolution   of   structure   dur ing   the  oxidation  of  zirconium  diboride-silicon  carbide  in  air  up  to C. J Eur Ceram Soc 2007;27(6):2495-501, http://dx.doi.org/10.1016/  1500     j.jeurceramsoc.2006.10.012.  3. Levine   S, Opila   E, Halbig M, Kiser   J,   Singh M,   Salem   J.   Eval uation   of   ultra-high  temperature  ceramics  2002;22(14-15):2757-67,   for   aeropropulsion   use.  J Eur Ceram   Soc   http://dx.doi.org/10.1016/  S0955-2219(02)00140-1.  4. Glass DE. Physical challenges an limitations confronting the use of UHTCs  on hypersonic vehicles. In: 17th AIAA international space planes and hyper sonic systems and technologies conference. 2011.  5. Sciti D,   Savino R,   Silvestroni   L. Aerothermal   behaviour   of   a   SiC  fibre-reinforced   ZrB2  sharp  component  2012;32(8):1837-45,   in   supersonic   regime.  J   Eur   Ceram   Soc   http://dx.doi.org/10.1016/  j.jeurceramsoc.2012.01.019.  6. Driemeyer PE, Haggerty RP, Kim D, Bell   JL, Apostolov ZD, Kriven  WM.   In   situ   studies   of   oxidation   of   ZrB2  and  ZrB2 -SiC  compos2010;30(11):2375-86,  ites   at   high   temperatures.   J   Eur   Ceram   Soc   http://dx.doi.org/10.1016/j.jeurceramsoc.2010.03.009.  7.   Jayaseelan DD, Zapata-Solvas E, Brown   P, Lee WE.   In   situ   forma tion   of   oxidation  resistant  refractory  coatings  on SiC-reinforced ZrB2 temperature ceramics. J Am Ceram Soc 2012;95(4):1247-54,  ultra high   http://dx.doi.org/10.1111/j.1551-2916.2011.05032.x.  8. Grigoriev ON, Galanov BA, Lavrenko VA, Panasyuk AD,   Ivanov SM,  Koroteev AV, Nickel KG. Oxidation of ZrB2 -SiC-ZrSi2 ceramics  in oxygen. J Eur Ceram Soc 2010;30(11):2397-405, http://dx.doi.org/10.1016/  j.jeurceramsoc.2010.03.016.  9. Hilmas G, Wuchina E. Unpublished.  10. Carney   CM. Oxidation   resistance   of   hafnium   diboride-silicon  car2009;44(20):5673-81,  bide   from   1400   to   2000     C.   J   Mater   Sci   http://dx.doi.org/10.1007/s10853-009-3799-7.  11. Chen D, Li W, Zhang X, Hu P, Han J, Hong C, Han W. Microstructural fea ture and thermal shock behavior of hot-pressed ZrB2 -SiC-ZrO2 composite. Mater Chem Phys 2009;116:348-52.  12. Gao D, Zhang Y, Xu C, Song Y, Shi X. Effect   of   ozone   adsorption  on  the oxidation behavior of ZrB2 -SiC ceramic composites. Corros Sci 2011;53:840-4.  13. Guo W-M, Zhang G-J. Oxidation  resistance  and  strength  ZrB2 -SiC ceramics. J Eur Ceram Soc 2010;30:2387-95. 14. Opila E, Levine S, Lorincz J. Oxidation of ZrB2 and HfB2 -based ultratemperature ceramics: effect of Ta additions. J Mater Sci 2004;39: high   retention   of  5969-77.  15. Wu H, Zhang W. Fabrication and properties of ZrB2 -SiC-BN machinable ceramics. J Eur Ceram Soc 2010;30:1035-42.  16. NIST Property Data Summaries, Sintered Silicon Carbide; 2001.  17. Thermophysical properties database of materials for light water reactors and  heavy water reactors, IAEA-TECDOC-1496; 2006.  18. Evert A, Tzou Y, Hasselman D, Raj R. The   effect of particle-size on  the  thermal-conductivity of Zns diamond composites. Acta Metal Mater 1992;40(1):123-9, http://dx.doi.org/10.1016/0956-7151(92)90205-S.  19. Gupta D.  Influence of  solute  segregation on grain boundary energy and self-diffusion. Metall Trans A 1977;8A:1431-8.                    \\x0c']"
},{
  "_id": 196,
  "PDF": "Oxidation-ablation behaviors of hafnium carbide-silicon carbonitride systems at 1500 and 2500 C.pdf",
  "Text": "[\"Ceramics International xxx (xxxx) xxx-xxx  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www .e lsev ie r .com / loca te /ce ram in t  Oxidation/ablation behaviors of hafnium carbide-silicon carbonitride systems at 1500 and 2500 C  Na Nia,c, Wei Haoa,b,c,∗, Tianyu Liud, Lei Zhoud, Fangwei Guob,c, Xiaofeng Zhaob,c,∗∗, Ping Xiaoa  a Key Lab of Education Ministry for Power Machinery and Engineering, School of Mechanical Engineering, Shanghai Jiao Tong University, Shanghai, 200240, China b Shanghai Key Laboratory of High-Temperature Materials and Precision Forming, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, China c Gas Turbine Research Institute, Shanghai Jiao Tong University, Shanghai, 200240, China d School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, 710072, China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Hafnium carbide Silicon carbonitride Oxidation Ablation Self-healing  The oxide scales of hafnium carbide (HfC) typically exhibit a porous structure after oxidation/ablation due to the release of gas oxidation products, which allows oxygen penetration to promote the rapid oxidation of the HfC matrices. Here, we report that the oxidation/ablation resistance of HfC was enhanced by the incorporation of amorphous silicon carbonitride (SiCN). HfC-SiCN ceramics with 10 vol % SiCN showed a significant improvement in the oxidation/ablation resistance compared with pure HfC. The HfC-10 vol % SiCN ceramic has a higher density with good mechanical properties. After being oxidized at 1500 °C for 2 h, a dense and homogeneous HfO2-HfSiO4 layer with low oxygen permeability is formed. The ablation resistance of the HfC-10 vol % SiCN ceramic is improved due to the formation of the triple-layer structure oxide with good thermal stability and mechanical scouring resistance. After ablation under an oxyacetylene flame for 60 s, the mass and linear ablation rates of HfC-10 vol % SiCN ceramic are −0.019 mg cm−2 s−1 and -0.156 μm s−1, respectively.  1.  Introduction  Hafnium carbide (HfC) is a promising candidate for thermal structural materials in hypersonic vehicles due to its good thermodynamic stability and high melting point (~3900 °C), together with its high hardness, Young's modulus, and thermal conductivity at ultra-high temperatures [1,2]. However, HfC is sensitive to oxygen-containing gases at ultra-high temperatures. Oxidation of HfC typically produces a porous HfO2 oxide layer due to the release of gas products, resulting in high oxygen permeability. This porous oxide layer can be easily scoured, and the characteristic limits the use of HfC in extreme environments [3,4]. To solve this problem, various oxidation-resistant additives, such as tantalum carbide (TaC) [2], tantalum (Ta) [5], silicon carbide (SiC) [6], oxides [7,8], silicides [9] and lanthanum boride (LaB6) [10] have been incorporated into HfC to induce the formation of a stable and dense oxide layer on the surface. Although a molten phase that heals the oxide layer at high temperatures can be formed after oxidation when TaC or  Ta is incorporated into HfC, the molten tantalum pentoxide (Ta2O5) layer that forms is not a promising oxidation-protective layer because phase transformations cause large changes in volume [2,5]. Yang et al. [6] reported the oxidation/ablation resistance of a HfC-SiC system, where a part of the silica in the formed Hf-Si-O compound layer was volatilized to leave pores and holes in the oxide layer during ablation. Introducing hafnia (HfO2) and yttria (Y2O3) into HfC was also investigated. However, these two materials have intrinsically high oxygen permeability, and a desirable molten phase that heals the porous oxide layer cannot be formed [8,11]. Incorporating silicides has been shown to induce the formation of a dense silicate layer that protects the matrices from oxidation and ablation, but an excess of molten silicates can deform oxide layers and cause holes to form during mechanical scouring [9]. When molybdenum disilicide (MoSi2) and tungsten disilicide (WSi2) are added, Moand W-oxides derived from the oxidation of their silicides can embrittle the protective oxide layer structure [12]. When LaB6 is incorporated into HfC, the release of gas is facilitated by the reduced melting point of the oxide scale; however, an increase in  ∗ Corresponding author. Shanghai Key Laboratory of High-Temperature Materials and Precision Forming, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, China. ∗∗ Corresponding author. Shanghai Key Laboratory of High-Temperature Materials and Precision Forming, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, China. E-mail addresses: haoweimaster@163.com (W. Hao), xiaofengzhao@sjtu.edu.cn (X. Zhao).  https://doi.org/10.1016/j.ceramint.2020.06.161 Received 23 March 2020; Received in revised form 9 June 2020; Accepted 15 June 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  Please cite this article as: Na Ni, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2020.06.161  \\x0c\", 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 1. XRD patterns (a), Raman spectra (b) and BSE-SEM images (c-f) of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN by SPS: (c) No SiCN; (d) 5 vol % SiCN; (e) 10 vol % SiCN; (f) 15 vol % SiCN.  2  \\x0c', \"N. Ni, et al.  the liquid boron oxide-lanthanum oxide (B2O3-La2O3) in the oxide scale lowers the resistance to mechanical scouring during ablation [10]. Therefore, to improve the oxidation/ablation resistance of HfC ceramics, the oxide layer that forms during the oxidation/ablation needs to have a dense and stable structure with a low oxygen permeability. Because HfC oxidation typically results in a porous HfO2 framework, the task is to tailor the secondary oxide phases to have the right combination of viscosity and strength to heal the porous structure, resist mechanical scouring, and allow an adequate release of gaseous products. SiCN is an ultra-high temperature material with excellent creep resistance, thermal stability, and oxidation/corrosion resistance at high temperatures [13-15]. In addition, SiCN oxidation can form a dense, self-healing silicate or silica layer with an oxygen diffusion rate (kp = 10−11 g2 cm4 s−1) that is lower than that of HfO2 (kp = 10−7 g2 cm4 s−1) [11,16,17]. In this work, different contents of SiCN were incorporated into HfC by spark plasma sintering (SPS), and the oxidation/ablation behaviors of the resulting HfC-SiCN ceramics were investigated. Mechanisms underlying the enhanced oxidation and ablation resistance that is observed with the addition of 10 vol % SiCN are discussed.  2. Experimental procedure  2.1. Spark plasma sintering (SPS)  Commercially available HfC powders (purity > 99.5%; impurities include ZrC < 0.46%, and O < 0.50%; the mean particle size is 100 nm; Shanghai Chao Wei Nano Technology Co. Ltd, China) were used as starting materials. SiCN having a particle size of 5-8 μm was obtained from Southwest Jiaotong University (China). The details of the powder synthesis were previously described [13,18,19]. HfC ceramics with 0, 5, 10, and 15 vol % SiCN were prepared by SPS using an SPS furnace (FTC HP D25, FCT Systeme GmbH, Rauenstein, Germany). Before starting the experiments, the SPS chamber pressure was lowered to 5 Pa, and argon was introduced to allow pyro-flushing for more stable temperature readings. HfC nanopowders were then poured into a 30 mm graphite die lined with 0.3-mm-thick graphite foil to maximize the electrical and thermal conduction between the punches and the die. The compacts were heated from 50 °C to 1100 °C at 100 °C/min and then heated from 1100 °C to 1850 °C at 50 °C/min. After a dwell time of 20 min at 50 MPa, the furnace was allowed to cool naturally to room temperature.  2.2. Characterization of  the as-sintered HfC-SiCN ceramics  2.2.1. Structure characterization The phase compositions of the as-sintered ceramics were characterized by X-ray diffraction (XRD; Ultima IV, Rigaku) taken with Cu Ka radiation (k = 0.15406 nm). Raman spectroscopy of the as-sintered samples was performed on a Raman microprobe spectroscopy (LabRAM HR, Horiba Jobin Yvon, France). A focused laser spot (Nd: YAG, 532 nm) with a diameter of approximately 3 μm was used. The microstructure and chemical composition of the samples were characterized by field emission scanning electron microscopy (FE-SEM; MIRA3LHM, TESCAN) equipped to perform energy dispersive X-ray spectroscopy (EDX) and scanning transmission electron microscopy (STEM; TALOS F200 × 200 kV, FEI) equipped with EDX. The density of the assintered samples was measured by the Archimedes method. Based on the densities of HfC (12.2 g cm−3) and SiCN (~2.2 g cm−3), the powders were mixed at varying volume fractions of SiCN (0, 5, 10, and 15 vol %) [11,20]. The theoretical densities of the corresponding samples are calculated to be 12.2 g cm−3, 11.7 g cm−3, 11.2 g cm−3,  Ceramics International xxx (xxxx) xxx-xxx  and 10.3 g cm−3, respectively. The grain size was estimated by the line intercept method using Image J software for image segmentation, and at least ten SEM images of each type of sample were used.  2.3. Mechanical  tests  The Young's modulus (E) and Vickers hardness (H) of the as-sintered samples were estimated from load-displacement curves produced by microindentation (Anton Paar, CPX MHT, Austria) with a diamond Vickers indenter. Fracture toughness was determined by the three-point single-edge notched beam method (SENB) using a universal mechanical testing machine (Zwick/Roell 155050-30 kN, Germany) with a 20 mm span and a crosshead speed of 0.05 mm min−1 during testing. The KIC values were calculated using Eq. (1) and Eq. (2) in accordance with the ASTM C-1421 standard.  (1)  (2)  where KIC is the fracture toughness (MPa m1/2); Y is the stress intensity factor coefficient; Pmax is the maximum force (N); S0 is the outer span (m); B is the width of the tested specimen (m); W is the height of the test specimen (m); a is the notch depth (m). Test bars of 3 mm × 4 mm × 24 mm were cut from the as-sintered HfC-SiCN ceramic pellets. The sample surfaces were abraded with diamond discs (50-100 μm) followed by diamond abrasive polishing down to a particle size of 0.5 μm. A U-groove was cut on the 3 mm × 35 mm surface using a 300-μm-thick diamond wheel, and the ratio of the notch depth to sample height (a/W) was controlled in the range of 0.35-0.6.  2.4. Oxidation tests  Thermogravimetric analysis (TGA) was performed using a thermal analysis system (STA449F3, NETZSCH, Germany) to investigate the oxidation behavior of the HfC-SiCN samples from room temperature to 1500 °C in air. Cubic samples (5 mm × 5 mm × 5 mm) were heated at a rate of 20 °C/min. Weight change vs. oxidation time was recorded in the thermogravimetric mode. Additionally, isothermal oxidation tests of the as-sintered HfC-SiCN samples were carried out at 1500 °C with natural convection of air. The thickness of the oxide layer was measured at room temperature after the isothermal oxidation tests using SEM and analyzed with Image J software. Tests were performed in triplicate, and the reported thicknesses are average values.  2.5. Ablation tests  The ablation behaviors of HfC and HfC-SiCN sample pellets (Φ30 × 5 mm) were evaluated under an oxyacetylene flame. The flame was parallel to the sample axis, and the distance between the torch nozzle and the sample surface was 10 mm. The heat flux during ablation was approximately 2400 kW m−2. The pressures of O2 and C2H2 were 0.4 and 0.095 MPa, respectively. The fluxes of O2 and C2H2 were 0.244 and 0.167 L s−1, respectively. After being tested under the flame for 60 s, the samples were naturally cooled to room temperature. During the test, an infrared thermometer indicated that the highest temperature of the central ablated surface reached approximately 2500 °C. The linear ablation rate (LAR) and mass ablation rate (MAR) of the samples were calculated according to Eq. (3) and Eq. (4),  3  =×××××KYPSBWaWaW103(/)2(1/),ICmax063/21/23/2==++YYaWaWaWaWaWaW(/)1.91095.1552(/)12.6880(/)19.5736(/)15.9377(/)5.1454(/),2345\\x0c\", \"N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Table 1 Density, grain size, and mechanical properties of the as-sintered HfC and HfC-SiCN ceramics.  Sample  Density (g cm−3)  Relative density (%)  Grain size (nm)  Hardness (GPa)  Young's modulus (GPa)  Fracture toughness (MPa m1/2)  HfC HfC-5 vol. % SiCN HfC-10 vol % SiCN HfC-15 vol % SiCN  11.5 11.3 10.9 10.3  94.3 95.6 96.3 96.8  680 ± 260 630 ± 180 600 ± 210 520 ± 200  15.8 ± 0.4 13.8 ± 0.2 13.0 ± 0.2 11.2 ± 0.3  286.7 ± 6.8 279.0 ± 6.6 273.5 ± 6.5 238.7 ± 6.9  4.3 ± 0.5 5.2 ± 0.4 7.2 ± 0.5 8.5 ± 0.5  in the elemental composition of the grains. Differences in grain orientation may also be a cause of the different contrasts seen in the BSESEM images. Combined with our previous investigation on these systems using TEM [19], we propose that the contrast is primarily due to small variations in the elemental composition of the grains. Some of dark contrast in HfC ceramic sintered without SiCN may be pores seen in the BSE-SEM images (Fig. 1c). In terms of the dark contrast at grain boundary areas, it is suggested to be associated with the formation of free carbon during the oxidation of HfC in SPS [21], which is based on the complementary TEM characterization of the same type of samples in our previous work. Notably, as the SiCN content increases, the grain size of HfC-SiCN ceramics gradually decreases from ~700 nm to ~520 nm (Table 1), and SiCN grains with dark contrast are homogeneously distributed in bulk HfC-SiCN ceramics, as shown in Fig. 1c-f. Densification mechanisms were discussed in our previous work [19] and are considered to be due to the synergistic effects of grain refinement; the formation of a solid solution of Si, O, and N; and segregation of carbon at grain boundaries. As SiCN content increases, Young's modulus and hardness decrease gradually (Table 1) because of the softer turbostratic carbon component of SiCN. The fracture toughness increases as SiCN content increases and reaches a maximum value of 8.5 ± 0.5 MPa m1/2 in samples containing 15 vol % SiCN, which is consistent with the results of our previous work [19].  3.2. Oxidation behavior of  the as-sintered HfC-SiCN ceramics  To investigate the oxidation behaviors of the as-sintered HfC and HfC-SiCN ceramics between 200 and 1500 °C, the oxidation resistance of the ceramics was characterized by TGA (Fig. 2). Without SiCN, the weight of the HfC ceramic increases by 1.78%; with 5 vol % SiCN, the weight of the HfC-SiCN ceramic increases by 3.49%. With 10 vol % SiCN, the HfC-SiCN exhibits excellent oxidation resistance within the tested temperature range. The maximum weight gain of the sample at 1500 °C is only 1.25%. With 15 vol % SiCN, the weight of the HfC-SiCN ceramic increases by 3.78%. The oxidation process of each sample can be divided into three phases, marked as A, B and C, as shown in Fig. 2. From 200 to 600 °C (phase A), the weight increases of all samples are similar. The increases are only slight as the ceramic samples react with oxygen slowly below 600 °C [11,22]. From 600 to 1300 °C (phase B), the weights of samples increase rapidly due to the enhanced oxidation. Silicate is not expected to have sufficient viscosity to seal the defects in the oxide layer at ~1300 °C [17]. From 1300 to 1500 °C (phase C), the weights of all samples increase. With HfC and HfC-10 vol % samples, oxidation accelerates. The oxidation decelerates in HfC-SiCN samples with 5 and 15 vol % SiCN. Severe oxidation in these two samples during phase B may promote the formation of more silicates, which are beneficial for retarding oxygen diffusion at higher temperatures, resulting in slow oxidation during phase C [17,23]. XRD patterns of the as-sintered HfC and HfC-SiCN ceramics after TGA at 200-1500 °C in air are shown in Fig. 3a. After TGA analyses, the  Fig. 2. wt change curves of the as-sintered HfC and HfC-SiCN ceramics in air with increasing temperature from 200 to 1500 °C at a heating rate of 20 °C/min.  respectively. The reported ablation rates were the average values of triplicate specimens.  (3)  (4)  where Rl refers to LAR (μm s−1) and Rm to MAR (mg cm−2 s−1); l0 is the thickness of the as-sintered samples (μm); l1 is the sample thickness after ablation (μm); m0 is the mass of the as-sintered samples (mg); m1 is the sample mass after ablation (mg); Δt is the ablating time (s); S is the surface area of samples (cm−2).  3. Results  3.1. Structure of  the as-sintered HfC-SiCN ceramics  The XRD patterns acquired from the HfC and HfC-SiCN ceramics sintered by SPS are shown in Fig. 1a. The indexed diffraction peaks demonstrate that all samples are primarily composed of a major HfC phase (PDF NO. 65-8747) and a minor HfO2 phase (PDF NO. 65-1142). These results suggest that the oxidation of HfC occurs during the SPS process. Raman spectra also confirmed that the as-sintered samples contain a free carbon phase derived from the addition of SiCN (1354 cm−1 D peak, 1585 cm−1 G peak), as shown in Fig. 1b [18]. Fig. 1c-f shows backscattered electron scanning electron microscopy (BSE-SEM) images of HfC ceramics sintered with or without SiCN. The samples consist of relatively uniform small grains (grain size: ~700 nm) with slightly different contrasts, which may be due to minor variations  4  =Rllt,l01=×RmmSt,m01\\x0c\", 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 3. XRD patterns (a, b) and SEM surface images (c-f1) of the as-sintered HfC and HfC-SiCN ceramics after TGA at 200-1500 °C in air: (c, c1) No SiCN; (d, d1) 5 vol % SiCN; (e, e1) 10 vol % SiCN; (f, f1) 15 vol % SiCN. Note that the surface oxide layer has been removed to investigate the crystal structure of the underlying nonoxidized HfC and HfC-SiCN samples in (b).  diffraction peaks of the cubic HfC phase disappear, and monoclinic HfO2 peaks are observed. As SiCN content increases, diffraction peaks of HfSiO4 are detected on the surface of the oxide layers in samples with 10 and 15 vol % SiCN, and the higher intensity of the associated diffraction peaks indicates that more HfSiO4 forms in the 15 vol % SiCN sample. These results suggest that the oxidation of HfC and HfC-SiCN samples up to 1500 °C can be described by the following reactions (Eqs. (5)-(7)) [3,6,22]:  (5)  (6)  (7)  Because residual carbon is also present in the HfC-SiCN samples, Eq. (8) should also occur.  (8)  XRD patterns of the underlying HfC and HfC-SiCN samples after TGA are shown in Fig. 3b (the surface oxide layer was removed by polishing). These results indicate that the crystal structure of the HfC matrix in each sample is still cubic (PDF NO. 65-8747) without any phase transformation after TGA. SEM surface images of HfC and HfCSiCN samples after TGA are shown in Fig. 3(c-f1). Without SiCN, large horizontal and vertical cracks are observed in the formed oxide layer (Fig. 3c). With 5 vol % SiCN, the number of cracks gradually decreases in the oxide layers (Fig. 3(d, d1)). Notably, when the SiCN content is 10 vol %, a denser oxide layer with fewer microcracks is obtained (Fig. 3(e, e1)). For the HfC-15 vol % SiCN sample, additional oxidation holes are observed in the oxide layers (Fig. 3(f, f1)). To investigate the effect of SiCN on the oxidation behavior of HfCSiCN ceramics in greater detail, isothermal oxidation tests were  5  ++HfC2OHfOCO.sgsg()2()2()2()+++SiCN3OSiOCONO.sglgg()2()2()2()2()+HfOSiOHfSiOsls2()2()4()+COCO.sgg()2()2()\\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 4. Oxidation behavior of the as-sintered HfC and HfC-SiCN ceramics at 1500 °C in air: thickness of the oxide layer during the first 1000 s of the oxidation.  (a)  thickness of  the oxide layer as a function of oxidation time,  (b)  performed at 1500 °C in air. Isothermal oxidation curves of HfC-SiCN ceramics with different SiCN contents are shown in Fig. 4a. Without SiCN, HfC exhibits poor oxidation resistance. After 4 h oxidation, the oxide thickness rapidly increases to ~1560 μm. With 5 vol % SiCN, the oxide thickness gradually decreases to ~420 μm after 8 h oxidation. Notably, HfC-10 vol % SiCN samples exhibit good oxidation resistance, and the thickness of the oxide layer is only ~200 μm after 8 h oxidation. However, with 15 vol % SiCN, the oxide thickness increases to ~900 μm after 8 h oxidation. At the initial stage of oxidation (~30 min) (Fig. 4b), the oxide thicknesses vary in a trend that is similar to the entire oxidation process. XRD patterns of the as-sintered HfC-SiCN ceramics after oxidation at 1500 °C in air are shown in Fig. 5. At 30 s, the oxidation products contain only monoclinic HfO2 in both the HfC and HfC-SiCN samples (Fig. 5b). At 30 min and 2 h, the oxidation products contain primarily tetragonal HfO2 instead of monoclinic HfO2 in the HfC ceramic samples. In the HfC-SiCN samples, the diffraction peaks can be indexed to HfSiO4 and monoclinic HfO2 (Fig. 5c, d). Overall, the addition of SiCN induces HfSiO4 formation and delays the monoclinic to tetragonal phase transformation during the oxidation process, which may be due to the suppression of phase transformation with the formation of silicate [24]. In the HfC-10 vol % SiCN samples, the peak intensity of monoclinic HfO2 increases with oxidation time and the transformation to tetragonal HfO2 occurs only after 8 h of oxidation (Fig. 5a). SEM surface images of the as-sintered HfC and HfC-SiCN ceramics after isothermal oxidation at 1500 °C in air for different oxidation times are shown in Fig. 6 (a-e1). The number and size of the cracks that form in the oxide layer decrease gradually as the SiCN content increases from 0 to 10 vol %. However, micropores and microcracks are generated in the oxide layer of samples with 15 vol % SiCN during oxidation (Fig. 6(c3, d3)). The evolution of defects in the oxide layer with oxidation time differs between samples with and without SiCN content. In samples without SiCN, the width of the formed cracks increases to ~5 μm as oxidation time increases from 10 s to 2 h (Fig. 6c and d). Similar behavior is observed for samples with 5 vol % SiCN (Fig. 6(b1d1)). However, for samples with 10 vol % (Fig. 6(b2-d2) and 15 vol % SiCN (Fig. 6(b3-d3)), the number of microcracks and micropores decreases as oxidation time increases from 10 s to 2 h. Notably, a dense  and homogeneous oxide layer is formed after 2 h of oxidation in samples with 10 vol % SiCN (Fig. 6d2). To investigate the phase distribution in the oxide layer, SEM-EDX analyses were performed on the surface of samples with 10 vol % SiCN after 30 min of oxidation. The oxide layer is composed of two types of grains with a dark or bright contrast (Fig. 6e2). The grains with the dark contrast (Spot 1 in Fig. 6e2) primarily contain Hf, Si, and O (Fig. 6e3). The grains that appear gray (Spot 2 in Fig. 6e4) contain only Hf and O (Fig. 6e4). These results suggest the presence of HfSiO4 and HfO2 phases and agree with the results of the XRD analyses. SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics after different oxidation times are shown in Fig. 7. In samples without SiCN, the thickness of the oxide layer rapidly increases from ~20 μm at 30 s to ~1300 μm at 2 h (Fig. 7a-d). Large vertical and horizontal cracks are observed in the oxide layer, interface and HfC matrix after 2 h of oxidation (Fig. 7e). With 5 vol % SiCN, microcracks are observed in the oxide layer, and their size increases with increasing oxidation time (Fig. 7(a1-d1)). With 10 vol % SiCN, a dense and homogeneous oxide layer of ~220 μm froms after 2 h of oxidation (Fig. 7(a2-d2)). With 15 vol % SiCN, the oxidation rate increases, and the oxide layer thickness reaches ~450 μm after 2 h of oxidation and induces microcracks in the oxide layer (Fig. 7(a3-d3)). The local area SEM-EDX analyses (Fig. 7c2) confirm that the content of Si gradually decreases from the HfC-SiCN matrix to the oxide layer, and that of O rapidly increases from the HfC-SiCN matrix to the oxide layer (Table 2). These results demonstrate that Si diffuses into the interface between the oxide layer and the matrix and participates in oxidation protection. Simultaneously, the diffusion of O2 in the dense composite oxide layer decreases.  3.3. Ablation behavior of  the as-sintered HfC-SiCN ceramics  The mass and linear ablation rates of the as-sintered HfC and HfCSiCN ceramics after ablation at 2500 °C for 60 s are shown in Fig. 8. During the ablation process, ceramic samples experience oxidation and mechanical scouring simultaneously; these processes cause mass gain/ thickness increase and mass loss/thickness reduction, respectively [25]. The mass and linear ablation rates reflect the net effect of oxidation and  6  \\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 5. XRD patterns of the as-sintered HfC and HfC-SiCN ceramics after oxidation at 1500 °C in air: (a) HfC-SiCN ceramic with 10 vol % SiCN after different oxidation times; (b) after a 30 s oxidation; (c) after a 30 min oxidation; (d) after a 2 h oxidation.  mechanical scouring. Without SiCN, the mass and linear ablation rates of HfC are 0.031 mg cm−2 s−1 and 0.733 μm s−1, respectively. As the content of SiCN increases, the mass and linear ablation rates gradually decrease. With 10 vol % SiCN, the mass and linear ablation rates are −0.019 mg cm−2 s−1 and -0.156 μm s−1, respectively. The slightly negative values indicate controllable mass gain and thickening, and suggest that the oxide layer forms stably, exhibits good thermal stability, and resists damage from mechanical scouring. As the SiCN content increases to 15 vol %, the samples experience large weight gain and become increasingly thick, which indicates that the oxidation of the sample during ablation is severe. Surface XRD patterns are shown in Fig. 9 and SEM surface images of the as-sintered HfC and HfC-SiCN ceramics after ablation tests are presented in Fig. 10. Regardless of the SiCN content, the sample  surfaces are composed of monoclinic HfO2. Without SiCN, some buckling and micropores are observed (Fig. 10a and a1). With 5 vol % SiCN, the number and size of microcracks increase (Fig. 10b and b1). Notably, with 10 vol % SiCN, a dense and homogeneous oxide layer is formed (Fig. 10c and c1). With 15 vol % SiCN, the density of holes on the sample surface increases (Fig. 10d and d1). SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics after ablation are shown in Fig. 11. A typical double-layer structure of HfO2, containing a dense upper layer and a porous bottom layer, is observed in the HfC sample (Fig. 11a and a1). Microcracks between the substrate and the oxide layer as well as the penetrating cracks in the oxide layer are seen in the ceramic sample with 5 vol % SiCN (Fig. 11b and b1). With 10 vol % SiCN, no cracks are present in the oxide layer, although the oxide layer is significantly thicker. This  7  \\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 6. SEM surface images (a-e1) of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN after isothermal oxidation at 1500 °C in air with different oxidation times, and SEM-EDX spectra (e2, e3) of the as-sintered HfC-SiCN ceramics with 10 vol % SiCN after 30 min oxidation.  oxide layer exhibits a unique triple-layer structure, including a dense upper layer, a porous intermediate layer and another dense bottom layer (Fig. 11c and c1). This oxide layer structure is different from those reported with HfC-based ceramics [2,9,26]. In the oxide layer of samples with 15 vol % SiCN, large holes are seen (Fig. 11d and d1), which is consistent with the surface observations. The oxide layer exhibits a uniform one-layer structure. The incorporation of 15 vol % SiCN induces the formation of excessive silica in the HfO2-SiO2 oxide layer (Fig. 12a, c). The segregation of silica is detected in the large holes (Fig. 12a, b).  SEM-EDX area analyses and elemental mapping were performed to investigate the composition of different oxide layers in HfC-10 vol % SiCN samples after ablation (Fig. 13). The dense upper layer contains Hf, C, Si, and O (“Area 1” in Fig. 13 and Table 3), which indicates that the oxide layer is composed of HfO2 and silica. From the cross-sectional analysis of the ablated sample, the formation of silica can also be inferred (Fig. 12). The amorphous nature of silica explains the absence of any Si-containing phase in the XRD patterns. The intermediate layer exhibits a porous microstructure containing Hf, C, and O (“Area 2” in Fig. 13 and Table 3). The dense bottom layer contains Hf, C, Si, and O  8  \\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 7. SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN after oxidation at 1500 °C in air with different oxidation times (inset of figure (e) represents the local magnification of (d)).  Table 2 Elemental compositions of different areas in the oxide layer of HfC-SiCN ceramics with 10 vol % SiCN after oxidation at 1500 °C in air for 30 min that are shown in Fig. 7c2.  EDX Area analyses  Area 1 Area 2 Area 3 Area 4  Elemental compositions  at. % Hf 17.81 16.60 17.32 6.62  C 37.24 39.38 42.37 77.12  Si 2.46 2.52 2.60 5.61  O 42.49 41.50 37.71 10.65  (“Area 3” in Fig. 13 and Table 3). The HfC matrix primarily contains Hf, C, and Si, as expected, with a small amount of O (“Area 4” in Fig. 13 and Table 3). These results are consistent with the elemental composition data of HfC-SiCN samples investigated in our previous work [19]. The porous intermediate and dense bottom layers in the HfC-10 vol % SiCN ceramic samples were characterized by STEM and STEM-EDX analyses. Fig. 14a and Fig. 14b show STEM bright-field (BF) and STEM high angle annular dark-field (HAADF) images, respectively. These images exhibit the dense and porous structure of the different oxide layers and are in good agreement with the SEM results. For the  intermediate porous layer (“Area d” in Fig. 14b), a large number of HfO2 grains are found to form a framework (Fig. 14(d1-d4) and Fig. 14h) due to the severe oxidation during ablation. Several HfC matrices (Fig. 14g) and SiOC grains (Fig. 14f) are observed in the HfO2 framework. In the bottom dense layer (“Area e” in Fig. 14b), bright grain, dark grain, and gray grain boundaries are observed. EDX mapping suggests that these are HfC, SiOC, and HfO2, respectively (Fig. 14(e1-e4)). The interfacial region between the SiOC and the HfC grains was analyzed by high-resolution transmission electron microscopy (HRTEM) (Fig. 14c1 and Fig. 14c2). Lattice fringes with spacings of approximately 0.31 and 0.27 nm are measured; these correspond to the d-spaces of ( 11) and (111) planes of monoclinic HfO2, respectively. A set of fringes in the adjacent SiOC grains is measured to be approximately 0.25 nm, corresponding to the d-space of the (002) plane of hexagonal SiC [19,27]. STEM-EDX mapping including the dark and bright areas in Fig. 14c1 and Fig. 14c2 is presented in Fig. 14(e-e4). It can be seen clearly that O is enriched at the grain boundary locations corresponding the dark areas in Fig. 14c1 and Fig. 14c2, as marked by the arrows in Fig. 14e4. This confirms the presence of HfO2 at grain boundaries, which is suggested by the lattice fringe analyses in Fig. 14c1 and Fig. 14c2. In addition, the dark contrast in the Z contrast dominated STEM HAADF image contains mainly Si, C and O (Fig. 14e).  9  1\\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 8. Ablation performance of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN addition at 2500 °C for 60 s: (a) mass ablation rate; (b) linear ablation rate.  oxidation provides channels for rapid oxygen diffusion (Figs. 6c and 7b). Microcracks were also observed in the oxidation of (Hf0.2Zr0·2Ta0·2Nb0·2Ti0.2)C [28]. We propose that these microcracks are generated in the formed oxide layer during thermal cycling due to the stress of thermal mismatch (coefficients of thermal expansion: HfC: 6.7 × 10−6 K−1; HfO2: 5.8 × 10−6 K−1; HfSiO4: 3.9 × 10−6 K−1) [11,29]. In addition, the observed phase transformation from monoclinic to tetragonal (Fig. 5) can also induce a volume change and contribute to the formation of microcracks (Fig. 6) [30]. Slower oxidation is achieved by increasing the SiCN content up to 10 vol %, where the formation of HfSiO4 decreases the volatilization of SiO2 (Eq. (6)). The HfO2 phase forms a framework in the oxide layer, where silicates can fill the HfO2 backbone and heal the defects [17,23]. In the HfC-10 vol % SiCN samples, the supply of Si leads to the formation of silicate and a dense and homogeneous HfO2-HfSiO4 composite oxide layer (Figs. 5 and 7c2 and Table 2) that inhibits the inward diffusion of oxygen and results in the best oxidation protection. In addition, the phase transformation of HfO2 is retarded by the addition of SiCN, which reduces the number of microcracks in the oxide layer. However, when the SiCN content increases to 15 vol %, more rapid oxidation occurs. Penetrating microcracks and microholes are observed in the oxide layer (Fig. 6d3 and Fig. 7c3). In this case, a larger amount of gases (CO2 and NO2) are produced through the reactions described by Eq. (5), Eq. (6), and Eq. (8). It appears that the amount of molten silicate is not sufficient to heal the holes when these gases are escaping from the oxide, as has been proposed previously [3,31]. Therefore, we propose that the generation of gas products during oxidation should be minimized. Another factor to consider is the thermal mismatch between the oxide layer and HfC matrix (coefficients of thermal expansion: HfC: 6.7 × 10−6 K−1; HfSiO4: 3.9 × 10−6 K−1) [11,29]. The thicker oxide layer formed after oxidation in the HfC-15 vol % SiCN samples may generate higher thermal stress and promote defect formation. The oxidation resistance of HfC-SiCN samples based on the weight gain from TGA is slightly different from that based on oxide layer thickness from isothermal oxidation. Although silicate is formed at ~1300 °C during a dynamic heating process, the shorter time at 1500 °C reduces the ability of the molten silicate to seal the oxide layer, which may have led to the large weight gain in the HfC-5 vol. % SiCN samples.  Fig. 9. Surface XRD patterns of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN after ablation.  The shape and intensity of the Si and C rich areas match with those of the Hf depletion areas characterized by the STEM-EDX elemental mapping analyses (Fig. 14(e1-e3)), confirming the presence of the SiOC grains with bright contrast in Fig. 14c1 and Fig. 14c2.  4. Discussion  4.1. Oxidation behavior  The incorporation of SiCN has a favorable effect on the oxidation resistance of HfC-SiCN ceramics at 1500 °C in air. Without SiCN, severe oxidation is caused by the high oxygen diffusion rate of pure HfO2 (kp = 10−7 g2 cm4 s−1), which is higher than that of silicates (kp = 10−11 g2 cm4 s−1) (Fig. 5). The formation of cracks during  10  \\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 10. SEM surface images of the as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN after ablation: (a, a1) No SiCN; (b, b1) 5 vol % SiCN; (c, c1) 10 vol % SiCN; (d, d1) 15 vol % SiCN.  Fig. 11. SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics after ablation: (a, a1) No SiCN; (b, b1) 5 vol % SiCN; (c, c1) 10 vol % SiCN; (d, d1) 15 vol % SiCN.  Fig. 12. SEM cross-sectional ablation.  4.2. Ablation resistance  The effects of SiCN addition on the ablation resistance of HfC-SiCN samples are similar to its effects on oxidation resistance. Without SiCN, the HfC samples exhibit the highest ablation rate (Fig. 8), and the  11  image (a) and elemental compositions of different areas  (b, c)  in the oxide layer  for HfC-SiCN ceramics with 15 vol % SiCN after  typical double-layer HfO2 structure is observed on the ablation surface (Fig. 11(a, a1)) [2,26]. Buckling and the formation of micropores in the upper HfO2 layer are due to volume expansion induced by rapid oxidation [28]. This upper layer cannot effectively protect the HfC matrix from rapid oxidation due to its high oxygen permeability [11].  \\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 13. SEM cross-sectional  image (a) and SEM-EDX element mapping analyses (b-e) of the as-sintered HfC-SiCN ceramics with 10 vol % SiCN after ablation.  Table 3 Elemental compositions of different areas in the oxide layer of HfC-SiCN ceramics with 10 vol % SiCN after ablation, as characterized by SEM-EDX.  EDX Area analyses  Area 1 Area 2 Area 3 Area 4  Elemental compositions  at. % Hf 14.00 11.17 12.90 23.78  C 57.46 71.30 45.90 67.72  Si 2.87 - 2.52 2.43  O 25.67 17.53 38.68 6.07  At the other extreme, the addition of 15 vol % SiCN also results in poor ablation resistance. The degradation is due to the formation of excessive silica in the HfO2-SiO2 oxide layer (Figs. 12 and 13), which softens the oxide layer and renders it less resistant to mechanical scouring during ablation. Higher SiCN content leads to the formation of higher level of CO2 and NO2 gas during ablation by the reactions described by Eq. (5), Eq. (6), and Eq. (8) and more CO, NO, and SiO being formed by the reaction described by Eq. (9). These products form because of the low partial pressure of oxygen during the ablation process, which causes large numbers of holes or pores to form in the oxide layer through volatilization [3,31]. The high thermal stress generated by a thick oxide layer may play a role in ablation resistance that is similar to its role in oxidation resistance, as discussed above.  (9)  Notably, HfC-10 vol % SiCN samples exhibit good ablation resistance at 2500 °C. This performance may be partially due to the mechanical properties associated with this sample. These samples exhibit a good balance between strength and toughness, which may contribute to  the resistance to cracking and mechanical scouring observed during ablation. More importantly, we propose that the unique triple-layer oxide structure that is formed in this sample (Fig. 11(c, c1) and Fig. 14) plays a critical role. The dense upper layer contains HfO2 and silica, which are formed by HfC oxidation and high-temperature sintering [2,32]. Furthermore, the phase diagram of the Si-HfO2 system suggests that only liquid SiO2 and monoclinic HfO2 are expected to exist at ~2000 °C [32]. Such a dense layer provides resistance to oxidation and mechanical scouring. A moderate SiCN content is beneficial for forming a silica-containing oxide layer that has an optimal viscosity for releasing gas products and sealing the oxide layer. The porous intermediate layer may be formed by the gas products that are retained in the oxide layer. The intermediate layer has an HfO2 framework containing the original HfC matrix and SiOC grains. Although this structure is porous, the matrix grains reinforce the oxide layer and resist mechanical scouring during ablation. In the dense bottom layer, a small amount of HfO2 is present at the boundary between HfC and SiOC. This result reveals that the slight oxidation in the interface between the bottom oxide layer and matrix occurs by preferential oxidation of the grain boundaries (Fig. 14 (c-c2) and Fig. 14e4), which is controlled by oxygen diffusion [28]. Therefore, the dense bottom layer is also expected to protect the HfC-SiCN matrix from oxidation [33,34]. Based on the finding of preferential oxidation of the grain boundaries, it is noted that the smallest grain size in the HfC15 vol % SiCN samples (Table 1) is expected to have the highest density of grain boundaries, which may contribute to its poor oxidation/ablation resistance. Further dedicated studies are needed in the future to clarify the effects of grain boundary oxidation.  12  +++2SiCN3O2SiO2CO2NO.sgggg()2()()()()\\x0c', 'N. Ni, et al.  Ceramics International xxx (xxxx) xxx-xxx  Fig. 14. TEM analysis of HfC-SiCN ceramics with 10 vol % SiCN after ablation: (a) STEM BF image; (b) STEM HAADF image; (c, c1, and c2) STEM HAADF image and HRTEM of a typical interfacial region between SiOC and HfC grains (the squared “Area c” in (e), as marked by the arrows in (e4)); (d, d1-d4) STEM HAADF image and corresponding STEM-EDX mapping of the squared “Area d” in (b); (e, e1-e4) STEM HAADF image and corresponding STEM-EDX mapping of the squared “Area e” in (b); (f, g, h) EDX spectra.  5. Conclusions  The oxidation and ablation behaviors of HfC-SiCN ceramics with different SiCN contents were investigated at 1500 and 2500 °C in air, respectively. The following conclusions can be drawn:  (i)  Incorporating SiCN has a beneficial effect on the oxidation and ablation resistance of HfC-based ceramics, and the best performance is achieved with 10 vol % SiCN. After ablation at 2500 °C, the mass and linear ablation rates of HfC-10 vol % SiCN samples are −0.019 mg cm−2 s−1and -0.156 μm s−1, respectively. (ii) The optimal oxidation resistance in HfC-10 vol % SiCN samples is due to the formation of a dense and homogeneous oxide layer, in which the HfO2 phase forms a framework, and silicates fill the spaces and heal the defects.  (iii)  Improving ablation resistance is associated with the formation of an oxide layer with a unique triple-layer structure in the HfC10 vol % SiCN samples. This structure has lower oxygen permeability and better mechanical strength for retarding oxidation and mechanical scouring during ablation.  Declaration of competing interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  13  \\x0c', 'Ceramics International xxx (xxxx) xxx-xxx  [24]  [22]  [21]  [15]  [20]  J.H. Niu, S.H. Meng, H. Jin, J.P. 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Ceram. Soc. 88 (2005) 2552-2558. J.F. Huang, L. Zhou, L.Y. Cao, H.B. Ouyang, W. Hao, C.Y. Li, J. Fei, Effect of the incorporation of SiC nanowire on mullite/SiC protective coating for carbon/carbon composites, Corrosion Sci. 107 (2016) 85-95. [32] D.W. Shi, R. Arroyave, Z.K. Liu, Thermodynamic modeling of the Hf-Si-O system, Calphad 30 (2006) 375-386. T. Kalvis, B.M. David, R. Rishi, Oxidation of polymer-derived HfSiCNO up to 1600°C, J. Am. Ceram. Soc. 96 (2013) 1278-1284. [34] M. Niu, Z.H. Zhao, L. Su, H.F. Gao, Z.X. Cai, H.J. Wang, Oxidation behavior of dense SiOC monolithics: the oxide scale development, Corrosion Sci. 163 (2020) 108235.  [27]  [28]  [31]  [33]  N. Ni, et al.  Acknowledgements  This work was supported by the National Natural Science Foundation of China (No. U19A2099 and 51902197) and the Shanghai Pujiang Program (No. 18PJ1406500). 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},{
  "_id": 197,
  "PDF": "Oxidation-ablation behaviors of hafnium carbide-silicon carbonitride systems at 1500 and 2500 C.pdf",
  "Text": "['Journal Pre-proof  Oxidation/ablation behaviors of hafnium carbide-silicon carbonitride systems at 1500 and 2500\\uffffC  Na Ni, Wei Hao, Tianyu Liu, Lei Zhou, Fangwei Guo, Xiaofeng Zhao, Ping Xiao  PII:  DOI:  S0272-8842(20)31820-4  https://doi.org/10.1016/j.ceramint.2020.06.161  Reference:  CERI 25588  To appear in:  Ceramics International  Received Date: 23 March 2020  Revised Date:  9 June 2020  Accepted Date: 15 June 2020  Please cite this article as: N. Ni, W. Hao, T. Liu, L. Zhou, F. Guo, X. Zhao, P. Xiao, Oxidation/ablation behaviors of hafnium carbide-silicon carbonitride systems at 1500 and 2500\\uffffC, Ceramics International (2020), doi: https://doi.org/10.1016/j.ceramint.2020.06.161.  This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.  © 2020 Published by Elsevier Ltd.  \\x0c', 'Oxidation/ablation behaviors of hafnium carbide-silicon carbonitride systems at   1500 and 2500 °C    Na Ni a, c, Wei Hao a, b, c,*, Tianyu Liu d, Lei Zhou d, Fangwei Guo b, c, Xiaofeng Zhao b,   c, *, Ping Xiao a    a Key Lab of Education Ministry for Power Machinery and Engineering, School of   Mechanical Engineering, Shanghai Jiao Tong University, Shanghai 200240, China   b Shanghai Key Laboratory of High-Temperature Materials and Precision Forming,   School of Materials Science and Engineering, Shanghai   Jiao Tong University,   Shanghai 200240, China   c Gas Turbine Research Institute, Shanghai Jiao Tong University, Shanghai 200240,   China    d School   of Materials Science   and Engineering, Northwestern Polytechnical   University, Xi ’an, 710072, China   Abstract:   The oxide scales of hafnium carbide (HfC) typically exhibit a porous structure after   oxidation/ablation due to the release of gas oxidation products, which allows oxygen   penetration to promote the rapid oxidation of the HfC matrices. Here, we report that   the oxidation/ablation resistance of HfC was enhanced by   the   incorporation of   amorphous silicon carbonitride (SiCN). HfC-SiCN ceramics with 10 vol. % SiCN   * Corresponding author: Tex./Fax: +86-21-54742561   E-mail address: haoweimaster@163.com (Wei Hao) and xiaofengzhao@sjtu.edu.cn   (Xiaofeng Zhao)                                                                   \\x0c', 'showed a significant improvement in the oxidation/ablation resistance compared with   pure HfC. The HfC-10 vol. % SiCN ceramic has a higher density with good   mechanical properties. After being oxidized at 1500   °C   for 2 h, a dense and   homogeneous HfO2-HfSiO4   layer with   low oxygen permeability   is formed. The   ablation resistance of the HfC-10 vol. % SiCN ceramic is improved due to the   formation of   the   triple-layer   structure oxide with good   thermal   stability   and   mechanical scouring resistance. After ablation under an oxyacetylene flame for 60 s,   the mass and linear ablation rates of HfC-10 vol. % SiCN ceramic are -0.019 mg cm-2   s-1 and -0.156 µm s-1, respectively.   Keywords: Hafnium carbide; Silicon carbonitride; Oxidation; Ablation; Self-healing   1.   Introduction   Hafnium carbide (HfC) is a promising candidate for thermal structural materials   in hypersonic vehicles due to its good thermodynamic stability and high melting point   (~3900   ºC),   together with   its high hardness, Young ’ s modulus,   and   thermal   conductivity   at ultra-high   temperatures   [1, 2]. However, HfC   is   sensitive   to   oxygen-containing gases at ultra-high   temperatures. Oxidation of HfC   typically   produces a porous HfO2 oxide layer due to the release of gas products, resulting in   high oxygen permeability. This porous oxide layer can be easily scoured, and the   characteristic limits the use of HfC in extreme environments [3-4].    To solve this problem, various oxidation-resistant additives, such as tantalum   carbide (TaC) [2], tantalum (Ta) [5], silicon carbide (SiC) [6], oxides [7-8], silicides   [9] and lanthanum boride (LaB6) [10] have been incorporated into HfC to induce the     \\x0c', 'formation of a stable and dense oxide layer on the surface. Although a molten phase   that heals the oxide layer at high temperatures can be formed after oxidation when   TaC or Ta is incorporated into HfC, the molten tantalum pentoxide (Ta2O5) layer that   forms is not a promising oxidation-protective layer because phase transformations   cause large changes in volume [2, 5]. Yang et al. [6] reported the oxidation/ablation   resistance of a HfC-SiC system, where a part of the silica in the formed Hf-Si -O   compound layer was volatilized to leave pores and holes in the oxide layer during   ablation. Introducing hafnia (HfO2) and yttria (Y2O3) into HfC was also investigated.   However, these two materials have intrinsically high oxygen permeability, and a   desirable molten phase that heals the porous oxide layer cannot be formed [8, 11].   Incorporating silicides has been shown to induce the formation of a dense silicate   layer that protects the matrices from oxidation and ablation, but an excess of molten   silicates can deform oxide layers and cause holes to form during mechanical scouring   [9]. When molybdenum disilicide (MoSi2) and tungsten disilicide (WSi2) are added,   Moand W-oxides derived from the oxidation of their silicides can embrittle the   protective oxide layer structure [12]. When LaB6 is incorporated into HfC, the release   of gas is facilitated by the reduced melting point of the oxide scale; however, an   increase in the liquid boron oxide-lanthanum oxide (B2O3-La2O3) in the oxide scale   lowers the resistance to mechanical scouring during ablation [10].    Therefore, to improve the oxidation/ablation resistance of HfC ceramics, the   oxide layer that forms during the oxidation/ablation needs to have a dense and stable   structure with a low oxygen permeability. Because HfC oxidation typically results in a   \\x0c', 'porous HfO2 framework, the task is to tailor the secondary oxide phases to have the   right combination of viscosity and strength   to heal   the porous structure,   resist   mechanical scouring, and allow an adequate release of gaseous products. SiCN is an   ultra-high temperature material with excellent creep resistance, thermal stability, and   oxidation/corrosion   resistance at high   temperatures   [13-15].   In addition, SiCN   oxidation can form a dense, self-healing silicate or silica   layer with an oxygen   diffusion rate (kp=10-11 g2 cm4 s-1) that is lower than that of HfO2 (kp=10-7 g2 cm4 s-1)   [11, 16, 17].   In this work, different contents of SiCN were incorporated into HfC by spark   plasma   sintering   (SPS),   and   the oxidation/ablation behaviors of   the   resulting   HfC-SiCN   ceramics were   investigated. Mechanisms   underlying   the   enhanced   oxidation and ablation resistance that is observed with the addition of 10 vol. % SiCN   are discussed.   2. Experimental Procedure   2.1 Spark plasma sintering (SPS)   Commercially available HfC powders (purity > 99.5 %; impurities include ZrC <   0.46 %, and O < 0.50 %; the mean particle size is 100 nm; Shanghai Chao Wei Nano   Technology Co. Ltd, China) were used as starting materials. SiCN having a particle   size of 5-8 µm was obtained from Southwest Jiaotong University (China). The details   of the powder synthesis were previously described [13, 18, 19]. HfC ceramics with 0,   5, 10, and 15 vol. % SiCN were prepared by SPS using an SPS furnace (FTC HP D25,   FCT Systeme GmbH, Rauenstein, Germany). Before starting the experiments, the       \\x0c', 'SPS chamber pressure was lowered to 5 Pa, and argon was introduced to allow   pyro-flushing for more stable temperature readings. HfC nanopowders were then   poured into a 30 mm graphite die lined with 0.3-mm-thick graphite foil to maximize   the electrical and thermal conduction between the punches and the die. The compacts   were heated from 50 °C to 1100 °C at 100 °C/min and   then heated from 1100 °C to   1850 °C at 50 °C/min. After a dwell time of 20 min   at 50 MPa, the furnace was   allowed to cool naturally to room temperature.   2.2 Characterization of the as-sintered HfC-SiCN ceramics   2.2.1 Structure characterization   The phase compositions of the as-sintered ceramics were characterized by X-ray   diffraction (XRD; Ultima IV, Rigaku) taken with Cu Ka radiation (k = 0.15406 nm).   Raman   spectroscopy of   the   as-sintered   samples was performed on   a Raman   microprobe spectroscopy (LabRAM HR, Horiba Jobin Yvon, France). A focused laser   spot (Nd: YAG, 532 nm) with a diameter of approximately 3 µm was used. The   microstructure and chemical composition of the samples were characterized by field   emission   scanning   electron microscopy   (FE-SEM; MIRA3-LHM,   TESCAN)   equipped   to perform energy dispersive X-ray spectroscopy (EDX) and scanning   transmission electron microscopy (STEM; TALOS F200X 200 kV, FEI) equipped   with EDX. The density of the as-sintered samples was measured by the Archimedes   method. Based on the densities of HfC (12.2 g cm-3) and SiCN (~2.2 g cm-3), the   powders were mixed at varying volume fractions of SiCN (0, 5, 10, and 15 vol. %)   [11, 20]. The theoretical densities of the corresponding samples are calculated to be       \\x0c', '12.2 g cm-3, 11.7 g cm-3, 11.2 g cm-3, and 10.3 g cm-3, respectively. The grain size was   estimated by the line intercept method using Image J software for image segmentation,   and at least ten SEM images of each type of sample were used.   2.2.2 Mechanical tests   The Young’s modulus ( E) and Vickers hardness (H) of the as-sintered samples   were estimated from load-displacement curves produced by microindentation (Anton   Paar, CPX MHT, Austria) with a diamond Vickers indenter. Fracture toughness was   determined by the three-point single-edge notched beam method (SENB) using a   universal mechanical testing machine (Zwick/Roell 155050-30 kN, Germany) with a   20 mm span and a crosshead speed of 0.05 mm min-1 during testing. The KIC values   were calculated using Eq. (1) and Eq. (2) in accordance with the ASTM C-1421   standard.   6  1/ 2  max  0  3/ 2  3 / 2  10  3  (  /  )  2  (1  /  )  IC  P  S  a W  K  Y  B W  a W   \\uf8ee \\uf8ef \\uf8f0  \\uf8f9 \\uf8ee \\uf8fa \\uf8ef \\uf8fb \\uf8f0  \\uf8f9 \\uf8fa \\uf8fb  · ·  ·  ·  =  ·   ,    (1)   (  )  (  )  (  )  (  )  (  )  2  3  4  5  (  /  )  1.91095.1552  /  12.6880  /  19.5736  /  15.9377  /  5.1454  /  Y  Y a W  a W  a W  a W  a W  a W  =  =  +   +   ,    (2)   where KIC   is   the fracture   toughness (MPa m1/2); Y   is   the stress   intensity factor   coefficient; Pmax is the maximum force (N); S0 is the outer span (m); B is the width of   the tested specimen (m); W is the height of the test specimen (m); a is the notch depth   (m).   Test bars of 3 mm × 4 mm × 24 mm were cut from the   as-sintered HfC-SiCN   ceramic pellets. The sample surfaces were abraded with diamond discs (50-100 µm)   followed by diamond abrasive polishing down   to a particle size of 0.5 µm. A   U-groove was cut on the 3 mm × 35 mm surface using   a 300-µm-thick diamond                             \\x0c', 'wheel, and the ratio of the notch depth to sample height (a/W) was controlled in the   range of 0.35-0.6.   2.2.3 Oxidation tests   Thermogravimetric analysis (TGA) was performed using a thermal analysis   system (STA449F3, NETZSCH, Germany) to investigate the oxidation behavior of   the HfC-SiCN samples from room temperature to 1500 °C in air. Cubic samples (5   mm×5 mm×5 mm) were heated at a rate of 20 °C/min. Weight ch ange vs. oxidation   time was recorded in the thermogravimetric mode. Additionally, isothermal oxidation   tests of the as-sintered HfC-SiCN samples were carried out at 1500 °C with natural   convection of air. The thickness of the oxide layer was measured at room temperature   after the isothermal oxidation tests using SEM and analyzed with Image J software.   Tests were performed in triplicate, and the reported thicknesses are average values.   2.2.4 Ablation tests   The ablation behaviors of HfC and HfC-SiCN sample pellets (Φ30 × 5 mm) were   evaluated under an oxyacetylene flame. The flame was parallel to the sample axis, and   the distance between the torch nozzle and the sample surface was 10 mm. The heat   flux during ablation was approximately 2400 kW m-2. The pressures of O2 and C2H2   were 0.4 and 0.095 MPa, respectively. The fluxes of O2 and C2H2 were 0.244 and   0.167 L s-1, respectively. After being tested under the flame for 60 s, the samples were   naturally cooled   to room   temperature. During   the   test, an   infrared   thermometer   indicated   that   the highest   temperature of   the central ablated   surface   reached   approximately 2500 °C. The linear ablation rate (LA R) and mass ablation rate (MAR)   \\x0c', 'of the samples were calculated according to Eq. (3) and Eq. (4), respectively. The   reported ablation rates were the average values of triplicate specimens.   l  0  =  R  l  l  1  ,   D  t  =  R  m  m  0  S  m  1  · D  t  ,    (3)    (4)   where Rl refers to LAR (µm s-1) and Rm to MAR (mg cm-2 s-1); l0 is the thickness of   the as-sintered samples (µm); l1 is the sample thickness after ablation (µm); m0 is the   mass of the as-sintered samples (mg); m1 is the sample mass after ablation (mg); Δt   is the ablating time (s); S is the surface area of samples (cm-2).   3. Results   3.1 Structure of the as-sintered HfC-SiCN ceramics   The XRD patterns acquired from the HfC and HfC-SiCN ceramics sintered by   SPS are shown in Fig. 1a. The indexed diffraction peaks demonstrate that all samples   are primarily composed of a major HfC phase (PDF NO. 65-8747) and a minor HfO2   phase (PDF NO. 65-1142). These results suggest that the oxidation of HfC occurs   during the SPS process. Raman spectra also confirmed that the as-sintered samples   contain a free carbon phase derived from the addition of SiCN (1354 cm-1 D peak,   1585 cm-1 G peak), as shown in Fig. 1b [18].   Fig.   1(c-f)   shows   backscattered   electron   scanning   electron microscopy   (BSE-SEM) images of HfC ceramics sintered with or without SiCN. The samples   consist of relatively uniform small grains (grain size: ~700 nm) with slightly different   contrasts, which may be due to minor variations in the elemental composition of the   grains. Differences in grain orientation may also be a cause of the different contrasts                                                                                                                               \\x0c', 'seen in the BSE-SEM images. Combined with our previous investigation on these   systems using TEM [19], we propose that the contrast is primarily due to small   variations in the elemental composition of the grains. Some of dark contrast in HfC   ceramic sintered without SiCN may be pores seen in the BSE-SEM images (Fig. 1c).   In terms of the dark contrast at grain boundary areas, it is suggested to be associated   with the formation of free carbon during the oxidation of HfC in SPS [21], which is   based on the complementary TEM characterization of the same type of samples in our   previous work.   Notably, as the SiCN content increases, the grain size of HfC-SiCN ceramics   gradually decreases from ~700 nm to ~520 nm (Table 1), and SiCN grains with dark   contrast are homogeneously distributed in bulk HfC-SiCN ceramics, as shown in Fig.   1(c-f). Densification mechanisms were discussed in our previous work [19] and are   considered to be due to the synergistic effects of grain refinement; the formation of a   solid solution of Si, O, and N; and segregation of carbon at grain boundaries. As SiCN   content   increases, Young ’s modulus and hardness dec rease gradually   (Table 1)   because of the softer turbostratic carbon component of SiCN. The fracture toughness   increases as SiCN content increases and reaches a maximum value of 8.5 ± 0.5 MPa   m1/2 in samples containing 15 vol. % SiCN, which is consistent with the results of our   previous work [19].   3.2 Oxidation behavior of the as-sintered HfC-SiCN ceramics   To investigate the oxidation behaviors of the as-sintered HfC and HfC-SiCN   ceramics between 200 and 1500 °C, the oxidation res  istance of the ceramics was     \\x0c', 'characterized by TGA (Fig. 2). Without SiCN,   the weight of   the HfC ceramic   increases by 1.78 %; with 5 vol. % SiCN, the weight of the HfC-SiCN ceramic   increases by 3.49 %. With 10 vol. % SiCN, the HfC-SiCN exhibits excellent oxidation   resistance within the tested temperature range. The maximum weight gain of the   sample at 1500 °C is only 1.25 %. With 15 vol. % Si CN, the weight of the HfC-SiCN   ceramic increases by 3.78 %. The oxidation process of each sample can be divided   into three phases, marked as A, B and C, as shown in Fig. 2. From 200 to 600 °C   (phase A), the weight increases of all samples are similar. The increases are only   slight as the ceramic samples react with oxygen slowly below 600 °C [11, 22]. From   600   to 1300   °C (phase B),   the weights of samples   in crease rapidly due   to   the   enhanced oxidation. Silicate is not expected to have sufficient viscosity to seal the   defects in the oxide layer at ~1300 °C [17]. From 1 300 to 1500 °C (phase C), the   weights of all samples increase. With HfC and HfC-10 vol. % samples, oxidation   accelerates. The oxidation decelerates in HfC-SiCN samples with 5 and 15 vol. %   SiCN. Severe oxidation in these two samples during phase B may promote the   formation of more silicates, which are beneficial for retarding oxygen diffusion at   higher temperatures, resulting in slow oxidation during phase C [17, 23].    XRD patterns of the as-sintered HfC and HfC-SiCN ceramics after TGA at   200-1500 °C in air are shown in Fig. 3a. After TGA analyses, the diffraction peaks of   the cubic HfC phase disappear, and monoclinic HfO2 peaks are observed. As SiCN   content increases, diffraction peaks of HfSiO4 are detected on the surface of the oxide   layers in samples with 10 and 15 vol. % SiCN, and the higher intensity of the   \\x0c', 'associated diffraction peaks indicates that more HfSiO4 forms in the 15 vol. % SiCN   sample. These results suggest that the oxidation of HfC and HfC-SiCN samples up to   1500 °C can be described by the following reactions   (Eq. (5-7)) [3, 6, 22]:   HfC  (  s  )  +  2O  2 (  g  )  fi  HfO  2 (  s  )  +  CO  2 (  g  )  SiCN  (  s  )  +  3O  2 (  g  )  fi  SiO  2 (  l  )  +  CO  2 (  g  )  +  NO  2 (  g  )  HfO  2 (  s  )  +  SiO  2 (  l  )  fi  HfSiO  4 (  s  )   (5)    (6)    (7)   Because residual carbon is also present in the HfC-SiCN samples, Eq. (8) should   also occur.   C  (  s  )   + O  2 (  g  )  fi  CO  2 (  g  )   (8)   XRD patterns of the underlying HfC and HfC-SiCN samples after TGA are   shown in Fig. 3b (the surface oxide layer was removed by polishing). These results   indicate that the crystal structure of the HfC matrix in each sample is still cubic (PDF   NO. 65-8747) without any phase transformation after TGA. SEM surface images of   HfC and HfC-SiCN samples after TGA are shown in Fig. 3(c-f1). Without SiCN,   large horizontal and vertical cracks are observed in the formed oxide layer (Fig. 3c).   With 5 vol. % SiCN, the number of cracks gradually decreases in the oxide layers (Fig.   3(d, d1)). Notably, when the SiCN content is 10 vol. %, a denser oxide layer with   fewer microcracks is obtained (Fig. 3(e, e1)). For the HfC-15 vol. % SiCN sample,   additional oxidation holes are observed in the oxide layers (Fig. 3(f, f1)).   To investigate the effect of SiCN on the oxidation behavior of HfC-SiCN ceramics   in greater detail,   isothermal oxidation   tests were performed at 1500   °C   in air.   Isothermal oxidation curves of HfC-SiCN ceramics with different SiCN contents are                                                                                                                                                                                                                                                                                           \\x0c', 'shown in Fig. 4a. Without SiCN, HfC exhibits poor oxidation resistance. After 4 h   oxidation, the oxide thickness rapidly increases to ~1560 µm. With 5 vol. % SiCN, the   oxide thickness gradually decreases to ~420 µm after 8 h oxidation. Notably, HfC-10   vol. % SiCN samples exhibit good oxidation resistance, and the thickness of the oxide   layer is only ~200 µm after 8 h oxidation. However, with 15 vol. % SiCN, the oxide   thickness increases to ~900 µm after 8 h oxidation. At the initial stage of oxidation   (~30 min) (Fig. 4b), the oxide thicknesses vary in a trend that is similar to the entire   oxidation process.   XRD patterns of the as-sintered HfC-SiCN ceramics after oxidation at 1500 °C   in air are shown in Fig. 5. At 30 s, the oxidation products contain only monoclinic   HfO2 in both the HfC and HfC-SiCN samples (Fig. 5b). At 30 min and 2 h, the   oxidation products contain primarily tetragonal HfO2 instead of monoclinic HfO2 in   the HfC ceramic samples. In the HfC-SiCN samples, the diffraction peaks can be   indexed to HfSiO4 and monoclinic HfO2 (Fig. 5(c, d)). Overall, the addition of SiCN   induces HfSiO4   formation   and   delays   the monoclinic   to   tetragonal   phase   transformation during the oxidation process, which may be due to the suppression of   phase transformation with the formation of silicate [24]. In the HfC-10 vol. % SiCN   samples, the peak intensity of monoclinic HfO2 increases with oxidation time and the   transformation to tetragonal HfO2 occurs only after 8 hours of oxidation (Fig. 5a).   SEM surface   images of   the as-sintered HfC and HfC-SiCN ceramics after   isothermal oxidation at 1500 °C in air for differen t oxidation times are shown in Fig. 6   (a-e1). The number and size of the cracks that form in the oxide layer decrease     \\x0c', 'gradually as the SiCN content increases from 0 to 10 vol. %. However, micropores   and microcracks are generated in the oxide layer of samples with 15 vol. % SiCN   during oxidation (Fig. 6(c3, d3)). The evolution of defects in the oxide layer with   oxidation time differs between samples with and without SiCN content. In samples   without SiCN, the width of the formed cracks increases to ~5 µm as oxidation time   increases from 10 s to 2 h (Fig. 6c and Fig. 6d). Similar behavior is observed for   samples with 5 vol. % SiCN (Fig. 6(b1-d1)). However, for samples with 10 vol. %   (Fig. 6(b2-d2) and 15 vol. % SiCN (Fig. 6(b3-d3)), the number of microcracks and   micropores decreases as oxidation time increases from 10 s to 2 h. Notably, a dense   and homogeneous oxide layer is formed after 2 h of oxidation in samples with 10 vol. %   SiCN (Fig. 6d2).   To investigate the phase distribution in the oxide layer, SEM-EDX analyses were   performed on the surface of samples with 10 vol. % SiCN after 30 min of oxidation.   The oxide layer is composed of two types of grains with a dark or bright contrast (Fig.   6e2). The grains with the dark contrast (Spot 1 in Fig. 6e2) primarily contain Hf, Si,   and O (Fig. 6e3). The grains that appear gray (Spot 2 in Fig. 6e4) contain only Hf and   O (Fig. 6e4). These results suggest the presence of HfSiO4 and HfO2 phases and agree   with the results of the XRD analyses.   SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics after   different oxidation times are shown in Fig. 7. In samples without SiCN, the thickness   of the oxide layer rapidly increases from ~20 µm at 30 s to ~1300 µm at 2 h (Fig.   7(a-d)). Large vertical and horizontal cracks are observed in the oxide layer, interface     \\x0c', 'and HfC matrix after 2 h of oxidation (Fig. 7e). With 5 vol. % SiCN, microcracks are   observed in the oxide layer, and their size increases with increasing oxidation time   (Fig. 7(a1-d1)). With 10 vol. % SiCN, a dense and homogeneous oxide layer of ~220   µm froms after 2 h of oxidation (Fig. 7(a2-d2)). With 15 vol. % SiCN, the oxidation   rate increases, and the oxide layer thickness reaches ~450 µm after 2 h of oxidation   and induces microcracks in the oxide layer (Fig. 7(a3-d3)).   The local area SEM-EDX analyses (Fig. 7c2) confirm that the content of Si   gradually decreases from the HfC-SiCN matrix to the oxide layer, and that of O   rapidly increases from the HfC-SiCN matrix to the oxide layer (Table 2). These   results demonstrate that Si diffuses into the interface between the oxide layer and the   matrix and participates in oxidation protection. Simultaneously, the diffusion of O2 in   the dense composite oxide layer decreases.   3.3 Ablation behavior of the as-sintered HfC-SiCN ceramics   The mass and   linear ablation rates of   the as-sintered HfC and HfC-SiCN   ceramics after ablation at 2500 °C for 60 s are sho wn in Fig. 8. During the ablation   process,   ceramic   samples   experience   oxidation   and   mechanical   scouring   simultaneously;   these processes   cause mass gain/thickness   increase   and mass   loss/thickness reduction, respectively [25]. The mass and linear ablation rates reflect   the net effect of oxidation and mechanical scouring. Without SiCN, the mass and   linear ablation rates of HfC are 0.031 mg cm-2 s-1 and 0.733 µm s-1, respectively. As   the content of SiCN increases, the mass and linear ablation rates gradually decrease.   With 10 vol. % SiCN, the mass and linear ablation rates are -0.019 mg cm-2 s-1 and         \\x0c', '-0.156 µm s-1, respectively. The slightly negative values indicate controllable mass   gain and thickening, and suggest that the oxide layer forms stably, exhibits good   thermal stability, and resists damage from mechanical scouring. As the SiCN content   increases   to 15 vol. %,   the samples experience   large weight gain and become   increasingly thick, which indicates that the oxidation of the sample during ablation is   severe.   Surface XRD patterns are shown in Fig. 9 and SEM surface images of the   as-sintered HfC and HfC-SiCN ceramics after ablation tests are presented in Fig. 10.   Regardless of the SiCN content, the sample surfaces are composed of monoclinic   HfO2. Without SiCN, some buckling and micropores are observed (Fig. 10a and Fig.   10a1). With 5 vol. % SiCN, the number and size of microcracks increase (Fig. 10b   and Fig. 10b1). Notably, with 10 vol. % SiCN, a dense and homogeneous oxide layer   is formed (Fig. 10c and Fig. 10c1). With 15 vol. % SiCN, the density of holes on the   sample surface increases (Fig. 10d and Fig. 10d1).   SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics after   ablation are shown in Fig. 11. A typical double-layer structure of HfO2, containing a   dense upper layer and a porous bottom layer, is observed in the HfC sample (Fig. 11a   and Fig. 11a1). Microcracks between the substrate and the oxide layer as well as the   penetrating cracks in the oxide layer are seen in the ceramic sample with 5 vol. %   SiCN (Fig. 11b and Fig. 11b1). With 10 vol. % SiCN, no cracks are present in the   oxide layer, although the oxide layer is significantly thicker. This oxide layer exhibits   a unique triple-layer structure, including a dense upper layer, a porous intermediate     \\x0c', 'layer and another dense bottom layer (Fig. 11c and Fig. 11c1). This oxide layer   structure is different from those reported with HfC-based ceramics [2, 9, 26]. In the   oxide layer of samples with 15 vol. % SiCN, large holes are seen (Fig. 11d and Fig.   11d1), which is consistent with the surface observations. The oxide layer exhibits a   uniform one-layer structure. The   incorporation of 15 vol. % SiCN   induces   the   formation of excessive silica   in   the HfO2-SiO2 oxide   layer (Fig. 12(a, c)). The   segregation of silica is detected in the large holes (Fig. 12(a, b)).   SEM-EDX area analyses and elemental mapping were performed to investigate   the composition of different oxide layers in HfC-10 vol. % SiCN samples after   ablation (Fig. 13). The dense upper layer contains Hf, C, Si, and O (“Area 1 ” in Fig.   13 and Table 3), which indicates that the oxide layer is composed of HfO2 and silica.   From the cross-sectional analysis of the ablated sample, the formation of silica can   also be inferred (Fig. 12). The amorphous nature of silica explains the absence of any   Si-containing phase in the XRD patterns. The intermediate layer exhibits a porous   microstructure containing Hf, C, and O (“Area 2 ” in   Fig. 13 and Table 3). The dense   bottom layer contains Hf, C, Si, and O (“Area 3 ” in   Fig. 13 and Table 3). The HfC   matrix primarily contains Hf, C, and Si, as expected, with a small amount of O (“Area   4 ” in Fig. 13 and Table 3). These results are consi  stent with the elemental composition   data of HfC-SiCN samples investigated in our previous work [19].   The porous intermediate and dense bottom layers in the HfC-10 vol. % SiCN   ceramic samples were characterized by STEM and STEM-EDX analyses. Fig. 14a   and Fig. 14b show STEM bright-field (BF) and STEM high angle annular dark-field   \\x0c', '(HAADF) images, respectively. These images exhibit the dense and porous structure   of the different oxide layers and are in good agreement with the SEM results. For the   intermediate porous layer (“Area d ” in Fig. 14b), a   large number of HfO2 grains are   found to form a framework (Fig. 14(d1-d4) and Fig. 14h) due to the severe oxidation   during ablation. Several HfC matrices (Fig. 14g) and SiOC grains (Fig. 14f) are   observed in the HfO2 framework. In the bottom dense layer (“Area e” in   Fig. 14b),   bright grain, dark grain, and gray grain boundaries are observed. EDX mapping   suggests that these are HfC, SiOC, and HfO2, respectively (Fig. 14(e1-e4)). The   interfacial   region   between   the SiOC   and   the HfC   grains was   analyzed   by   high-resolution   transmission electron microscopy (HRTEM) (Fig. 14c1 and Fig.   14c2). Lattice fringes with spacings of approximately 0.31 and 0.27 nm are measured;   these correspond to the d-spaces of (111) and (111) planes of monoclinic HfO2,   respectively. A set of   fringes   in   the adjacent SiOC grains   is measured   to be   approximately 0.25 nm, corresponding to the d-space of the (002) plane of hexagonal   SiC [19, 27]. STEM-EDX mapping including the dark and bright areas in Fig. 14c1   and Fig. 14c2 is presented in Fig. 14(e-e4). It can be seen clearly that O is enriched at   the grain boundary locations corresponding the dark areas in Fig. 14c1 and Fig. 14c2,   as marked by the arrows in Fig. 14e4. This confirms the presence of HfO2 at grain   boundaries, which is suggested by the lattice fringe analyses in Fig. 14c1 and Fig.   14c2. In addition, the dark contrast in the Z contrast dominated STEM HAADF image   contains mainly Si, C and O (Fig. 14e). The shape and intensity of the Si and C rich   areas match with those of the Hf depletion areas characterized by the STEM-EDX   \\x0c', 'elemental mapping analyses (Fig. 14(e1-e3)), confirming the presence of the SiOC   grains with bright contrast in Fig. 14c1 and Fig. 14c2.   4. Discussion   4.1 Oxidation behavior   The incorporation of SiCN has a favorable effect on the oxidation resistance of   HfC-SiCN ceramics at 1500 °C in air. Without SiCN,   severe oxidation is caused by   the high oxygen diffusion rate of pure HfO2 (kp = 10-7 g2 cm4 s-1), which is higher than   that of silicates (kp = 10-11 g2 cm4 s-1) (Fig. 5). The formation of cracks during   oxidation provides channels   for   rapid oxygen diffusion   (Fig. 6c and Fig. 7b).   Microcracks were also observed in the oxidation of (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)C [28].   We propose that these microcracks are generated in the formed oxide layer during   thermal cycling due   to   the stress of   thermal mismatch (coefficients of   thermal   expansion: HfC: 6.7 × 10  -6 K-1; HfO2: 5.8 × 10  -6 K-1; HfSiO4: 3.9 × 10  -6 K-1) [11, 29].   In addition, the observed phase transformation from monoclinic to tetragonal (Fig. 5)   can also induce a volume change and contribute to the formation of microcracks (Fig.   6) [30].    Slower oxidation is achieved by increasing the SiCN content up to 10 vol. %,   where the formation of HfSiO4 decreases the volatilization of SiO2 (Eq. (6)). The   HfO2 phase forms a framework in the oxide layer, where silicates can fill the HfO2   backbone and heal the defects [17, 23]. In the HfC-10 vol. % SiCN samples, the   supply of Si   leads   to   the formation of silicate and a dense and homogeneous   HfO2-HfSiO4 composite oxide layer (Fig. 5, Fig. 7c2 and Table 2) that inhibits the   \\x0c', 'inward diffusion of oxygen and results in the best oxidation protection. In addition,   the phase transformation of HfO2 is retarded by the addition of SiCN, which reduces   the number of microcracks in the oxide layer.   However, when the SiCN content increases to 15 vol. %, more rapid oxidation   occurs. Penetrating microcracks and microholes are observed in the oxide layer (Fig.   6d3 and Fig. 7c3). In this case, a larger amount of gases (CO2 and NO2) are produced   through the reactions described by Eq. (5), Eq. (6), and Eq. (8). It appears that the   amount of molten silicate is not sufficient to heal the holes when these gases are   escaping from the oxide, as has been proposed previously [3, 31]. Therefore, we   propose that the generation of gas products during oxidation should be minimized.   Another factor to consider is the thermal mismatch between the oxide layer and HfC   matrix (coefficients of thermal expansion: HfC: 6.7 × 10  -6 K-1; HfSiO4: 3.9 × 10  -6 K-1)   [11, 29]. The thicker oxide layer formed after oxidation in the HfC-15 vol. % SiCN   samples may generate higher thermal stress and promote defect formation.   The oxidation resistance of HfC-SiCN samples based on the weight gain from   TGA is slightly different from that based on oxide layer thickness from isothermal   oxidation. Although silicate is formed at ~1300 °C during a dynamic heating process,   the shorter time at 1500 °C reduces the ability of   the molten silicate to seal the oxide   layer, which may have led to the large weight gain in the HfC-5 vol. % SiCN samples.   4.2 Ablation resistance   The effects of SiCN addition on the ablation resistance of HfC-SiCN samples are   similar to its effects on oxidation resistance. Without SiCN, the HfC samples exhibit   \\x0c', 'the highest ablation rate (Fig. 8), and the typical double-layer HfO2 structure is   observed on the ablation surface (Fig. 11(a, a1)) [2, 26]. Buckling and the formation   of micropores in the upper HfO2 layer are due to volume expansion induced by rapid   oxidation [28]. This upper layer cannot effectively protect the HfC matrix from rapid   oxidation due to its high oxygen permeability [11].   At the other extreme, the addition of 15 vol. % SiCN also results in poor ablation   resistance. The degradation   is due   to   the   formation of excessive silica   in   the   HfO2-SiO2 oxide layer (Fig. 12 and Fig. 13), which softens the oxide layer and   renders it less resistant to mechanical scouring during ablation. Higher SiCN content   leads to the formation of higher level of CO2 and NO2 gas during ablation by the   reactions described by Eq. (5), Eq. (6), and Eq. (8) and more CO, NO, and SiO being   formed by the reaction described by Eq. (9). These products form because of the low   partial pressure of oxygen during the ablation process, which causes large numbers of   holes or pores to form in the oxide layer through volatilization [3, 31]. The high   thermal stress generated by a thick oxide layer may play a role in ablation resistance   that is similar to its role in oxidation resistance, as discussed above.   2SiCN    + 3O  (  s  )  2 (  g  )  fi  2SiO   (  g  )   + 2CO   (  g  )   + 2NO   (  g  )   (9)   Notably, HfC-10 vol. % SiCN samples exhibit good ablation resistance at   2500   °C. This performance may be partially due   to   t he mechanical properties   associated with this sample. These samples exhibit a good balance between strength   and toughness, which may contribute to the resistance to cracking and mechanical   scouring observed during ablation. More importantly, we propose that the unique                                               \\x0c', 'triple-layer oxide structure that is formed in this sample (Fig. 11(c, c1) and Fig. 14)   plays a critical role.    The dense upper layer contains HfO2 and silica, which are formed by HfC   oxidation and high-temperature sintering [2, 32]. Furthermore, the phase diagram of   the Si-HfO2 system suggests that only liquid SiO2 and monoclinic HfO2 are expected   to exist at ~2000 °C [32]. Such a dense layer provi des resistance to oxidation and   mechanical   scouring. A moderate SiCN   content   is beneficial   for   forming   a   silica-containing oxide layer that has an optimal viscosity for releasing gas products   and sealing the oxide layer.   The porous intermediate layer may be formed by the gas products that are   retained in the oxide layer. The intermediate layer has an HfO2 framework containing   the original HfC matrix and SiOC grains. Although this structure is porous, the matrix   grains reinforce the oxide layer and resist mechanical scouring during ablation. In the   dense bottom layer, a small amount of HfO2 is present at the boundary between HfC   and SiOC. This result reveals that the slight oxidation in the interface between the   bottom oxide layer and matrix occurs by preferential oxidation of the grain boundaries   (Fig. 14   (c-c2) and Fig. 14e4), which   is controlled by oxygen diffusion   [28].   Therefore, the dense bottom layer is also expected to protect the HfC-SiCN matrix   from oxidation [33, 34]. Based on the finding of preferential oxidation of the grain   boundaries, it is noted that the smallest grain size in the HfC-15 vol. % SiCN samples   (Table 1) is expected to have the highest density of grain boundaries, which may   \\x0c', 'contribute   to   its poor oxidation/ablation resistance. Further dedicated studies are   needed in the future to clarify the effects of grain boundary oxidation.   5. Conclusions   The oxidation and ablation behaviors of HfC-SiCN ceramics with different SiCN   contents were investigated at 1500 and 2500 °C in a ir, respectively. The following   conclusions can be drawn:   (i) Incorporating SiCN has a beneficial effect on the oxidation and ablation   resistance of HfC-based ceramics, and the best performance is achieved with 10 vol. %   SiCN. After ablation at 2500 °C, the mass and linea r ablation rates of HfC-10 vol. %   SiCN samples are -0.019 mg cm-2 s-1and -0.156 µm s-1, respectively.   (ii) The optimal oxidation resistance in HfC-10 vol. % SiCN samples is due to   the formation of a dense and homogeneous oxide layer, in which the HfO2 phase   forms a framework, and silicates fill the spaces and heal the defects.   (iii) Improving ablation resistance is associated with the formation of an oxide   layer with a unique triple-layer structure in the HfC-10 vol. % SiCN samples. This   structure has lower oxygen permeability and better mechanical strength for retarding   oxidation and mechanical scouring during ablation.   Acknowledgements   This work was supported by the National Natural Science Foundation of China   (No. U19A2099   and   51902197)   and   the   Shanghai   Pujiang   Program   (No.   18PJ1406500). Ablation   tests were performed   in   the Carbon/Carbon Composites   Research Center, School of Materials Science and Engineering, Northwestern       \\x0c', 'Polytechnical University.   References   [1] J.C. Ren, Y.L. Zhang, P.F. Zhang, T. Li, J.H. Li, Y. Yang, Ablation resistance of   HfC coating reinforced by HfC nanowires in cyclic ablation environment, J. Eur.   Ceram. Soc. 37 (2017) 2759 -2768.   [2] C. Zhang, B. Boesl, A. Agarwal, Oxidation resistance of tantalum carbide-hafnium   carbide solid solutions under the extreme conditions of a plasma jet, Ceram. Int.   43 (2017) 14798 -14806.   [3] M.D. Tong, Q.G. Fu, L. Zhou, T. Feng, H.J. Li, T. Li, K. Li, Ablation behavior of a   novel HfC-SiC gradient coating fabricated by a facile one-step chemical vapor   co-deposition, J. Eur. Ceram. 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Fei, Effect of   the   incorporation of SiC nanowire on mullite/SiC protective coating   for   carbon/carbon composites, Corros. Sci. 107 (2016) 85 -95.   [32] D.W. Shi, R. Arroyave, Z.K. Liu, Thermodynamic modeling of the Hf-Si-O   system, CALPHAD. 30 (2006) 375 -386.   [33] T. Kalvis, B.M. David, R. Rishi, Oxidation of Polymer-Derived HfSiCNO up to   1600 °C, J. Am. Ceram. Soc. 96 (2013) 1278 -1284.   [34] M. Niu, Z.H. Zhao, L. Su, H.F. Gao, Z.X. Cai, H.J. Wang, Oxidation behavior of   dense SiOC monolithics: the oxide scale development, Corros. Sci. 163 (2020)   108235.                     \\x0c', 'Figure Captions:   Fig. 1 XRD patterns (a), Raman spectra (b) and BSE-SEM images (c-f) of the   as-sintered HfC and HfC-SiCN ceramics with different contents of SiCN by SPS: (c)   No SiCN; (d) 5 vol. % SiCN; (e) 10 vol. % SiCN; (f) 15 vol. % SiCN.   Fig. 2 Weight change curves of the as-sintered HfC and HfC-SiCN ceramics in air   with increasing temperature from 200 to 1500 °C at a heating rate of 20 °C/min.   Fig. 3 XRD patterns (a, b) and SEM surface images (c-f1) of the as-sintered HfC and   HfC-SiCN ceramics after TGA at 200-1500 °C in air:   (c, c1) No SiCN; (d, d1) 5 vol. %   SiCN; (e, e1) 10 vol. % SiCN; (f, f1) 15 vol. % SiCN. Note that the surface oxide   layer has been   removed   to   investigate   the crystal   structure of   the underlying   non-oxidized HfC and HfC-SiCN samples in (b).                         \\x0c', 'Fig. 4 Oxidation behavior of the as-sintered HfC and HfC-SiCN ceramics at 1500 °C   in air: (a) thickness of the oxide layer as a function of oxidation time, (b) thickness of   the oxide layer during the first 1000 s of the oxidation.   Fig. 5 XRD patterns of the as-sintered HfC and HfC-SiCN ceramics after oxidation at   1500 °C in air: (a) HfC-SiCN ceramic with 10 vol. % SiCN after different oxidation   times; (b) after a 30 s oxidation; (c) after a 30 min oxidation; (d) after a 2 h oxidation.   Fig. 6 SEM surface images (a-e1) of the as-sintered HfC and HfC-SiCN ceramics   with different contents of SiCN after isothermal oxidation at 1500 °C in air with   different oxidation times, and SEM-EDX spectra (e2, e3) of the as-sintered HfC-SiCN   ceramics with 10 vol. % SiCN after 30 min oxidation.   Fig. 7 SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics   with different contents of SiCN after oxidation at 1500   °C in air with different   oxidation times (inset of figure (e) represents the local magnification of (d)).   Fig. 8 Ablation performance of the as-sintered HfC and HfC-SiCN ceramics with   different contents of SiCN addition at 2500 °C for   60 s: (a) mass ablation rate; (b)   linear ablation rate.   Fig. 9 Surface XRD patterns of the as-sintered HfC and HfC-SiCN ceramics with   different contents of SiCN after ablation.   Fig. 10 SEM surface images of the as-sintered HfC and HfC-SiCN ceramics with   different contents of SiCN after ablation: (a, a1) No SiCN; (b, b1) 5 vol. % SiCN; (c,   c1) 10 vol. % SiCN; (d, d1) 15 vol. % SiCN.   \\x0c', 'Fig. 11 SEM cross-sectional images of the as-sintered HfC and HfC-SiCN ceramics   after ablation: (a, a1) No SiCN; (b, b1) 5 vol. % SiCN; (c, c1) 10 vol. % SiCN; (d, d1)   15 vol. % SiCN.   Fig. 12 SEM cross-sectional image (a) and elemental compositions of different areas   (b, c) in the oxide layer for HfC-SiCN ceramics with 15 vol. % SiCN after ablation.   Fig. 13 SEM cross-sectional image (a) and SEM-EDX element mapping analyses (b-e)   of the as-sintered HfC-SiCN ceramics with 10 vol. % SiCN after ablation.   Fig. 14 TEM analysis of HfC-SiCN ceramics with 10 vol. % SiCN after ablation: (a)   STEM BF image; (b) STEM HAADF image; (c, c1, and c2) STEM HAADF image   and HRTEM of a typical interfacial region between SiOC and HfC grains (the squared   “Area c ” in (e), as marked by the arrows in (e4));   (d, d1-d4) STEM HAADF image   and corresponding STEM-EDX mapping of the squared “ Area d ” in (b); (e, e1-e4)   STEM HAADF image and corresponding STEM-EDX mapping of the squared “Area   e” in (b); (f, g, h) EDX spectra.   Table 1 Density, grain size, and mechanical properties of the as-sintered HfC and   HfC-SiCN ceramics.   Table 2 Elemental compositions of different areas in the oxide layer of HfC-SiCN   ceramics with 10 vol. % SiCN after oxidation at 1500 °C in air for 30 min that are   shown in Fig. 7c2.   Table 3 Elemental compositions of different areas in the oxide layer of HfC-SiCN   ceramics with 10 vol. % SiCN after ablation, as characterized by SEM-EDX.     \\x0c', 'Table 1 Density, grain size, and mechanical properties of the as-sintered HfC and   HfC-SiCN ceramics.   Sample   Density   Relative   Grain size   Hardness   (g cm-3)   density (%)   (nm)   (GPa)   Young ’s   Fracture   modulus   toughness   (GPa)   (MPa m1/2)   HfC   11.5   94.3   680±260   15.8±0.4   286.7±6.8   4.3±0.5   HfC-5 vol. % SiCN 11.3   95.6   630±180   13.8±0.2   279.0±6.6   5.2±0.4   HfC-10   vol.   %  10.9   96.3   600±210   13.0±0.2   273.5±6.5   7.2±0.5   SiCN   HfC-15   vol.   %  10.3   96.8   520±200   11.2±0.3   238.7±6.9   8.5±0.5   SiCN           \\x0c', 'Table 2 Elemental compositions of different areas in the oxide layer of HfC-SiCN   ceramics with 10 vol. % SiCN after oxidation at 1500 °C in air for 30 min that are   shown in Fig. 7c2.   EDX    Area analyses   Area 1   Area 2   Area 3   Area 4   Elemental compositions   at. %   Hf   C   Si   O   17.81   37.24   2.46   42.49   16.60   39.38   2.52   41.50   17.32   42.37   2.60   37.71   6.62   77.12   5.61   10.65     \\x0c', 'Table 3 Elemental compositions of different areas in the oxide layer of HfC-SiCN   ceramics with 10 vol. % SiCN after ablation, as characterized by SEM-EDX.   EDX    Area analyses   Elemental compositions   at. %   Hf   C   Si   O   Area 1   14.00   57.46   2.87   25.67   Area 2   11.17   71.30   —   Area 3   12.90   45.90   2.52   17.53   38.68   Area 4   23.78   67.72   2.43   6.07     \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', ' \\x0c', 'Declaration of interests   ☒ The authors declare that they have no known competing financial interests or personal relationships   that could have appeared to influence the work reported in this paper.   ☐The authors declare the following financial interests/personal relationships which may be considered   as potential competing interests:    We declare that we have no known competing financial interests or personal relationships that could   have appeared to influence the work reported in this paper.                                                       Na Ni, Wei Hao, Tianyu Liu, Lei Zhou, Fangwei Guo, Xiaofeng Zhao, Ping Xiao                                                                                                                                                                   Jun. 09, 2020                                                                                                     \\x0c']"
},{
  "_id": 198,
  "PDF": "Oxidation-based materials selection for 2000°C + hypersonic aerosurfaces Theoretical considerations and historical experience.pdf",
  "Text": "['ULTRA-HIGH TEMPERATURE CERAMICS  J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5887 - 5904  Oxidation-based materials selection for 2000 C + hypersonic aerosurfaces: Theoretical considerations and historical experience  M . M . O P E K A , I . G . T A LM Y , J . A . Z A Y K O S K I Naval Surface Warfare Center Carderock Division, West Bethesda, MD 20817, USA E-mail: opekaMM@nsweed.navy.mil  Hypersonic flight involves extremely high velocities and gas temperatures with the attendant necessity for thermal protection systems (TPS). New high temperature materials are needed for these TPS systems. A systematic investigation of the thermodynamics of potential materials revealed that low oxidation rate materials, which form pure scales of SiO2 , Al2O3 , Cr2O3 , or BeO, cannot be utilized at temperatures of 1800 C (and above) due to disruptively high vapor pressures which arise at the interface of the base material and the scale. Vapor pressure considerations provide significant insight into the relatively good oxidation resistance of ZrB2 and HfB2 -based materials at 2000 C and above. These materials form multi-oxide scales composed of a refractory crystalline oxide (skeleton) and a glass component, and this compositional approach is recommended for further development. The oxidation resistance of ZrB2 -SiC and other non-oxide materials is improved, to at least 1600 C, by compositional modifications which promote immiscibility in the glass component of the scale. Other candidate materials forming high temperature oxides, such as rare earth compounds, are largely unexplored for high temperature applications and may be attractive candidates for hypersonic TPS materials. C(cid:2) 2004 Kluwer Academic Publishers  1.  Introduction     The 21st century has ushered in a new, exciting era of hypersonic flight. Hypersonic flight vehicles include sub-orbital and earth-to-orbit vehicles for rapid global and space access missions. A common aspect of these future systems is the need for new high-temperature materials. Hypersonic vehicles with sharp aerosurfaces, such as engine cowl inlets, wing leading edges (LEs), and nosecaps, have projected needs for 2000 to 2400 C materials which must operate in air and be re-usable. At this time, there are few, if any, off-the-shelf materials to meet these future hypersonic thermal protection system (TPS) needs. State-of-the-Art high temperature materials include carbon-carbon composites (CC) and silicon carbide-based (SiC) composites, such as C-SiC and SiC-SiC. Ultra-High-Temperature Ceramics (UHTCs), such as Zr(Hf)B2 -SiC, are being developed but are less mature at this time. Carbon-carbon composites have very high temperature structural capabilities but are not oxidation-resistant. Coatings have been and are being developed for oxidation-resistance, but cyclic life capabilities are modest due to the difficulties of managing the thermal expansion coefficient (CTE) mismatch between the C-C composite and the coating systems. The SiC-based composites exhibit oxidation resistance up to 1600 C in hypersonic environments, but thermal cycling lifetimes are also modest due to CTEmismatch-induced matrix cracking which        allows direct oxidation of the carbon fiber reinforcement. The UHTCs, based on the diborides of zirconium and hafnium, have exhibited relatively good oxidation resistance above 1600 C. The oxidation mechanisms of these materials are not well understood. Recent Navy efforts to understand UHTC oxidation mechanisms, and to develop new, highly oxidation-resistant 2000 C materials, are presented here. This paper describes prior development of ultra-high-temperature, oxidation-resistant materials; thermodynamics and kinetics principles related to oxidation; theoretical aspects of the oxidation of UHTC materials; and experimental results associated with compositional variations of UHTC materials.     2. Developmental history  Distinct lines of research have contributed significantly to our current understanding of oxidationresistant ultra-high temperature materials: coating systems for refractory metals and subsequent development of oxidation-resistant intermetallic compounds, oxidation-resistant graphite compositions, and the development of boride-based UHTCs. The structural usefulness of refractory metals, and their lack of high temperature oxidation-resistance, motivated the pursuit of oxidation-resistant coatings. Considerable research was conducted, especially in the  0022-2461  C(cid:2) 2004 Kluwer Academic Publishers  5887  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS        1960s, and a number of texts summarizing the developments are available [1-5]. Although a broad range of materials was investigated, a significant proportion of the work was based on compositions containing silicon (Si), aluminum (Al), and chromium (Cr). Packer [6] summarized research on silicides, and an important conclusion is the significance of low pressure and high temperature environments on limiting the life of such materials and coatings. Perkins and Packer [7] identified the maximum temperature capability of MoSi2 coatings as 1800 C in atmospheric pressure (hypersonic) environments. Recent research on the oxidation of intermetallics, especially emphasizing aluminides for gas turbine applications, has been compiled by Grobstein and Doychak [8]. Oxidation-resistant graphite compositions were developed in parallel with refractory metal coatings in the 1960s [9, 10]. One of the most important compositions, designated “grade JTA” graphite [9], was optimized for oxidation resistance at 2000 C. It used additions of ZrB2 and Si (balance carbon) at an approximately 50 weight percent basis. Further optimization using transition metal additions (e.g., niobium) were found to improve oxidation performance at high temperatures, but with the penalty of poorer performance at lower temperatures. Krivoshein and coworkers [11] reported that Nb additions (10 wt%) improved oxidation performance of ZrB2 -SiC modified graphite, but that V additions at the same level provided maximum improvement. Significant research was reported on the refractory boride compounds beginning in the late 1940s with crystal structure [12] and melting point [13] measurements. An initial survey of the oxidation resistance of transition metal diborides up to 1500 C revealed that the group IVb compounds were the most resistant [14]. A survey of oxidation resistance of the diborides of Hf, Zr, Ti, Ta, and Nb from 1200 to 2200 C (inductively heated samples in flowing He-O2 mixtures) also revealed that HfB2 was the most oxidation resistant, followed by ZrB2 . The temperature dependence of the oxidation data for both compounds indicated significant rate changes at the respective metal oxide phase transition temperatures [15]. Oxidation testing in flowing He-O2 mixtures with H2O (at 613 Pa) exhibited a five-fold increase in oxidation rate of HfB2 at 933 C versus the nominal dry He-O2 . Similar oxidation measurements at 1487 C showed no rate difference [16]. Additional oxidation studies on ZrB2 and HfB2 demonstrated that metal-rich compositions (e.g., HfB1.7 ) oxidized at lower rates (by up to a factor of 50) versus boron-rich compositions (e.g., HfB2.12 ) [16]. Numerous investigations to improve the oxidation resistance of ZrB2 and HfB2 have been reported [17-20]. Compositions with 5 to 50 vol% SiC were investigated for both ZrB2 and HfB2 over a wide range of test temperatures and pressures; 20 vol% compositions were judged optimal for hypersonic vehicles in a series of efforts supported by the US Air Force [16, 19, 21-23]. Additions of C (5, 10, 15, 20, 30, and 50 vol%) improved thermal stress resistance, but were detrimental to oxidation resistance at all proportions. Additions of              5888     Cr (10 mol%), Al (20 mol%), and Ta (30 mol%) were found to be detrimental to oxidation resistance. An addition of 4 vol% of a Hf-20 at.%Ta alloy had no effect on the oxidation properties, although the metal phase was converted to the carbides during the hot-press fab50 mol% HfB2 + 50 mol% HfB composition exhibited rication process. Excess Hf metal additions to produce a rapid, preferential oxidation of the HfB phase. Additions of silicon to substitute on the boron sub-lattice yielded a HfB2 + “HfSi” skeletal phase which also exhibited rapid, preferential oxidation. (It is noted here that “HfSi” was identified by X-ray diffraction; other Hf-Si or Zr-Si second phases are possible, but have not been explored for oxidation response.) Additions of SiB6 (10 and 20 vol%) were found to increase oxidation resistance, but were not superior to SiC additions. Other systematic studies of additions into ZrB2 and or HfB2 have been conducted. Shaffer [24] evaluated the oxidation resistance of ZrB2 with additions of the disilicides of Ta, Nb, W, Mo, Zr, Mo0.5Ta0.5 , and Mo0.8Ta0.2 , as well as Zr5Si3 . The additive amounts were not specified, however, and only the conclusion was stated that MoSi2 was “unquestionably the best.” Additional oxidation experiments with varying proportions of MoSi2 (1 to 20 mol%) were conducted at 1950 C and revealed + 10 mol%MoSi2 composition was marketed by the the optimum composition to be 10 mol%. The ZrB2 Carborundum Company (US) under the Trade name “Boride Z”. Pastor and Meyer [25] evaluated the oxidation resistance of ZrB2 with additions of MSi2 or M5Si3 , where M is a transition metal Zr, Ta, Cr, Mo, or W. On the basis of scale thickness measurements after oxidation testing for up to 100 h at 1200 and 1400 + 15 wt%CrSi2 composition was found to be the most C, the ZrB2 oxidation resistant. Lavrenko and coworkers [26] reported that a ZrB2 + 50 wt%ZrSi2 composition was more oxidation resistant than MoSi2 and WSi2 , and could be used up to 1700 C. However, since oxidation data only up to 1200 C are reported, it is not clear how the conclusion is supported. The oxidation kinetics mechanism(s) of the diboridebased materials are only partly understood despite significant research. Oxidation kinetics measurements are typically based on weight change or scale thickness changes with time upon exposure to a known temperature and oxidizing atmosphere. However, weight change and scale thickness measurements are confounded by simultaneous oxidation and vaporization (of BOx vapor species) processes. Total oxygen consumption measurements (per unit area of sample) have been utilized to overcome this limitation [27]. Initial oxidation studies were conducted in 1955 on porous ZrB2 samples from 649 to 1315 C [28]. The oxidation kinetics were found to be parabolic, the rates increased with oxygen partial pressure, and the presence of H2O also increased the oxidation rate. Berkowitz-Mattuck measured total oxygen consumption for ZrB2 over a higher temperature range (1200-1530 C) and a lower oxygen partial pressure ( P o2 ) range (1070 to 5200 Pa) in helium (He) at 1.01 × 105 Pa total pressure [27]. Parabolic rate kinetics were                 \\x0c', '                  also observed, as were modest increases in oxidation rates with increasing P o2 . From metallographic examination of tested samples, it was concluded that oxidation proceeded by inward diffusion of oxygen, and it was suggested that oxygen diffusion through ZrO2 was the rate controlling step. Kuriakose and Margrave measured weight changes for ZrB2 over the temperature range of 945-1256 C and also reported parabolic oxidation kinetics [27, 29]. At 1056 C they observed that the parabolic rate con1.36 × 104 to 9.92 × 105 Pa at 1.01 × 105 Pa total presstant increased directly proportional to P o2 (range was sure with balance He). Berkowitz-Mattuck extended the oxidation kinetics studies of ZrB2 to understand the change in P o2 dependence with temperature [30, 31]. The P o2 dependence was confirmed at a test temperature of 927 C, but no dependence was found at 1557 C. Additional testing also revealed a significant change in the activation energy at 1057 C, changing from 20 kcal/mole below this temperature to 70 kcal/mole above it. Abrupt changes in the oxidation rate kinetics were also observed at the temperatures corresponding to the monoclinic to tetragonal oxide phase transitions for both ZrB2 and HfB2 . Other oxidation studies have been conducted on the oxidation kinetics of ZrB2 , HfB2 , and their respective SiC-modified compositions [32-40]. Oxygen diffusion through the B2O3 liquid phase was identified as the rate limiting step associated with oxidation of the pure diborides up to approximately 1200 C. Above this temperature, the increased oxidation rates were attributed to oxygen transport through the ZrO2 or HfO2 phase. The addition of SiC was found to significantly increase the temperature range of the glass as the primary oxygen barrier. A two layer scale was observed to form with HfO2 inner and SiO2 outer components. The reduction in oxidation rate was observed above 1350 C. Below this temperature, SiC inclusions are found in the HfO2 scale since the SiC particles do not oxidize significantly to generate the SiO2 glass component [34]. Low temperature oxidation studies have also been conducted for these borides. Preferential oxidation of C at an oxygen pressure of 1.3 × the Zr or Hf at 500 −3 Pa has been reported. Boron inclusions, which co10 alesced into layers, were observed in the oxide scale. Solution of oxygen into the diboride lattice was also reported [36]. Changes in the oxidation mechanism were noted at approximately 500 C [37]. In addition to the diborides, other materials were investigated for potential hypersonic applications [20, 41-44]. Materials based on ZrC and HfC were extensively studied, but were found to oxidize (nonprotectively) below 1800 C, which eliminated them from consideration for the temperature cycling hypersonic applications. Additions of SiC did not solve the rapid oxidation at low-temperatures. Hafniumtantalum alloys (e.g., Hf-20 at.%Ta) were found to exhibit good oxidation behavior, but were limited by the relatively low melting point of 2000 C at this composition. Iridium coatings on graphite were also evaluated, but were judged costly and not sufficiently refractory due to the Ir-C eutectic at 2296 C. The viscosity of                    ULTRA-HIGH TEMPERATURE CERAMICS     SiO2 was significantly increased by the addition of W powder in 10 and 20 vol% additions. These materials exhibited increased sensitivity to thermal stress failure and deformed by viscous flow into blunt shapes. Since hypersonic applications typically require the high temperature materials to retain sharp radii for leading edges and propulsion inlets, this shape change was unacceptable. By the early-1970s, the ZrB2 and HfB2 -based materials were identified to be the most promising for hypersonic applications with cyclic exposure from ambient temperature up to 2700 C [18, 41, 42]. Prior materials development for hypersonic applications does not include significant emphasis on oxide materials. They have not been pursued for these applications due to the demanding structural and thermostructural requirements of such systems, and low thermal shock resistance of oxides in general. It must be asked whether the optimized choice of hypersonic materials should be oxide or non-oxide materials. Oxide materials are, at best, intrinsically resistant to oxidation. However, oxide-oxide composites for 2000 C usage do not currently exist, and current and developmental oxide composites for aircraft applications cannot be used to that temperature. Such ultra-high-temperature oxide composites will likely be very costly due to the need to develop new creep-resistant reinforcements and suitable fiber-matrix interface materials to evade brittle fracture behavior. In addition, oxides and oxide composites incur significant design penalties due to their relatively high CTE and stiffness, and low thermal conductivity. Such a new 2000 C oxide composite would have to be developed for dedicated hypersonic application at very high cost. However, the payoff of intrinsic oxidation resistance requires a continuing look into oxide materials for these applications. The materials selection process presented here addresses the optimization of non-oxide ceramic compositions for high temperature hypersonic applications.        3. Oxidation-thermodynamics and kinetics  The selection of new oxidation-resistant materials is based on chemical thermodynamics and kinetics. Chemical thermodynamics is a powerful tool for identifying the equilibrium phases associated with the oxidation kinetics process(es). Chemical thermodynamics can be seen as providing the boundary conditions for understanding the oxidation kinetics processes. The thermodynamics-based calculations provided in this paper are based on the following relation:  \\x01G = \\x01G    + R · T · ln(Q )  (1)  where \\x01G is the change in Gibbs Free Energy (superscript refers to standard state) associated with a given chemical reaction, R is the ideal gas constant, T is the reaction temperature, and Q is the activity quotient [45, 46]. At the condition of chemical equilibrium, \\x01G is zero and Equation 1 reduces to:    = − R · T · ln(k )  \\x01G  (2)  5889  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  or,    = −2.303 · R · T · log(k )  \\x01G  (3)  where k is the reaction equilibrium constant. Gibbs Free Energies of Formation (FEOF) at standard state (\\x01G o f ) have been measured, and/or estimated, and tabulated as a function of temperature for many compounds of interest [47-50].  3.1. Thermodynamics and condensed phase equilibria  A stability diagram (Fig. 1) showing regions of metal and condensed metal oxide equilibria using Equation 3 was constructed for the elements Be, Y, Hf, Ta, W, Re, and Ir. These elements were selected as the most refractory representatives of their respective columns or groups in the periodic chart of elements. The diagram provides the metal-metal oxide equilibria as functions of temperature and oxygen pressure, P o2 . Each metal is stable below its respective metal-metal oxide equilibrium line, while the condensed oxide is stable above the equilibrium line. This diagram allows the generalized hypersonic environment to be directly compared with the metal-metal oxide stability regions. The FEOF data for BeO, HfO2 , Ta2O5 , and Re2O7 were taken from Schick [47], for WO3 from “JANAF” [48] , for Y2O3 from Pankratz [49], and for IrO2 from Knacke [50]. As necessary, FEOF were extrapolated linearly from the highest temperature data available. For Y2O3 , WO3 , Re2O7 , and IrO2 , the data were extrapolated above 1723, 2723, 361, and 1027 C, respectively, to compare with the hypersonic vehicle environment. Multiple condensed oxides exist and were considered for W (WO2 and WO3 ) and Re (ReO2 , ReO3 , and Re2O7 ). For both metals, the suboxides (WO2 , ReO2 , and ReO3 ) decompose below 1500 C and were neglected. The hypersonic environment envelope divides the diagram of Fig. 1 into three distinct groups of elements. The first group includes noble metals, such as iridium (Ir), for which the metal is the equilibrium condensed phase in the hypersonic environment. Since a condensed oxide does not form, mass loss from Irx O y        Figure 1 Metal-metal oxide condensed phase equilibrium diagram.  5890     vapor species must be considered. Oxidation mass loss rates for Ir are extremely low [51-53]. Other possible noble elements include rhodium, platinum, palladium, osmium, ruthenium, gold, and silver, which are not shown. Osmium and Ru exhibit very high mass loss rates in oxidizing environments [53]. Of the remaining elements, only iridium, and possibly rhodium, have melting points high enough to be considered for hypersonics applications. They should be considered for development but are not discussed further in this paper. The second element group shown in Fig. 1 includes those elements which form condensed oxides in the hypersonic environment, but the oxide melting point is below the 2000 C upper use temperature. These elements include tantalum, tungsten, rhenium, and by analogy, aluminum, titanium, molybdenum and all other group Vb (vanadium, niobium) and VIIb (manganese, technetium) elements. Since the oxides of these elements have relatively low melting points and melt viscosities, aerodynamic shear forces would remove them quickly and high ablation rates would be observed. In addition to the low oxide melting point limitation, these metals oxidize and experience high mass loss rates (from oxidative vaporization) at modest to high temperatures. Oxidation rates for Ta and W (as well as for Nb and Mo) have been reviewed by Kofstad [54]. High oxidation rates for Re in air are reported in the 300 to 1500 C temperature range [55, 56]. The third group, hafnium (Hf), yttrium (Y), and beryllium (Be), includes those elements which exist only as solid, condensed oxides throughout the hypersonic environment. Although only these three elements are shown, the result may be generalized to include zirconium, chromium, silicon, scandium, and the rare earth metals. Silicon is also included in this group because of the high viscosity of the molten oxide. Since these elements have melting points near to or lower than 2000 C, they have to be used as refractory compounds (silicides, borides, carbides, nitrides, etc.). Thus, initial candidate hypersonic materials include noble, refractory metals (Ir and/or Rh), and compounds (based on Zr, Hf, Be, Si, Y, Sc, and rare earths) which form refractory oxides.        3.2. Oxidation kinetics  While equilibrium thermodynamics provides an important foundation for materials selection, a second foundation is oxidation kinetics. For hypersonic applications, materials with a low oxidation rate are required. The need for higher temperature gas turbine materials has motivated significant investigations into very low oxidation rate materials. As shown in Fig. 2, Zr and Hf exhibit relatively high oxidation rates, while the most slowly oxidizing materials at high temperatures are those which form scales composed of pure SiO2 , Al2O3 , Cr2O3 , or BeO [57-64]. Above approximately 1100 C, SiO2 scale-forming materials (SiC, Si3N4 , MoSi2 , and possibly other silicides) have the lowest known oxidation rates [65, 66]. Beryllides (e.g., Ta2Be17 and ZrBe13 ) have very low oxidation rates up to 1250 C, but the rates increase rapidly with temperature [62, 63].        \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 2 Parabolic oxidation rate constants for various metals and com pounds.  3.3. Thermodynamics and vapor phase equilibria  Equation 3 may also be employed to construct diagrams of the vapor pressures of metal and metal oxide gaseous species as a function of oxidant pressure. Such a vapor phase diagram [67, 68], or volatility diagram, for the silicon-oxygen (Si-O) system at 2227 C is shown in Fig. 3. The vertical dashed line, at an oxygen pressure ( P o2 ) of 3.0×10 −5 Pa, is the equilibrium line defined by the reaction:     Si(c, l) + O2 (g) → SiO2 (l)  (4)  and separates the diagram into the two regions in which condensed Si and SiO2 exist. Unit activity is assumed for both Si and SiO2 , which is typical for this type diagram. Other lines show the vapor pressures of the SiO gaseous species (Si, Si2 , Si3 , SiO, SiO2 ) as a function of P o2 . The diagram may be seen as a schematic of the vapor pressure changes from the external surface of the SiO2 scale (right side of diagram) on oxidized Si, to the interior of the Si (left side of the diagram). It should be noted that the vapor pressures are highest at the Si-SiO2 interface (SiO vapor) and not at the exterior surface of the SiO2 . This high interfacial vapor pressure limits the use temperature of SiO2 -forming mate T A B L E I Calculated vapor pressures for oxides  Dominant  log P MOx  log PMOx  Condensed  Oxide  Melt temp  vapor species  at 1727     C  at 2227     C  for P = 10 Temp (deg C) −4 Pa  Temp (deg C) for P = 105 Pa  Material  oxide  type  (deg C)  (at Interface)  (Pa)  (Pa)  O2 Barrier matls Be  BeO  Crystalline  2550  Be  2.85  4.43  820  2495  Si  SiO2 Al2O3 Cr2O3 B2O3  Amorphous  1725  SiO  4.5  6.06  750  1865  Al  Crystalline  2040  Al, Al2O Cr  2.87  4.42  800  2490  Cr  Crystalline  2300  2.14  3.9  970  2690  B  Amorphous  450  B2O3 , B2O2  4.11  5.88  800  1950  Scale structures  Be  BeO  Crystalline  2550  Be  2.85  4.43  820  2495  Sc  Sc2O3 Y2O3 ZrO2 HfO2 Ta2O5  Crystalline  2400  Sc  2.02  3.8  950  2750  Y  Crystalline  2430  Y  −2.18 0.626 −2.58 −1.92  2.63  1120  3330  Zr  Crystalline  2700  ZrO  0.926  1520  3640  Hf  Crystalline  2800  HfO  0.506  1570  3670  Ta  Crystalline  1890  TaO2  0.926  1490  3730  Figure 3 Volatility diagram for Si-SiO2 system at 2227     C.  rials. The SiO2 scale is continuously ruptured when the interfacial (SiO) pressure exceeds ambient pressure and the scale loses its protective capability. At this condition, the SiO2 scale is consumed via the reaction:  Si(l) + SiO2 (l) → 2SiO(g)  (5)  The SiO(g) diffuses outward from the base of the porous SiO2 (l) scale and reacts with O2 (g) to form SiO2 (l), either within the scale or as smoke outside the scale [68]. If the SiO2 scale is entirely consumed (or if not formed initially), oxidation proceeds via the reaction:  Si(l) + 1/2O2 (g) → SiO(g)  (6)  The conditions represented by reactions (5) and (6) yield rapid oxidation of the Si. Wagner first showed that the transition between the rapid, (active) oxidation of Si(c,l) which occurs via reaction (6) and the slow (passive) oxidation via reaction (4) was governed primarily by the thermodynamics of the system [69]. Gulbransen includes reaction (5) in his description of active oxidation [68], although other researchers have restricted the term to describe the effect of reaction (6) only [70, 71]. The maximum in PSiO at the Si-SiO2 interface uniformly increases with temperature. The calculated SiO exceeds 1.01 × 105 Pa interfacial vapor pressure at 1865 C as shown in Table I.     5891  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS                    Gulbransen et al. [72] confirmed the active-passive transition boundary line for elemental Si in the 1100 to iments to air at 1.01 × 105 Pa yields the onset of active 1300 C temperature range. Extrapolating these experoxidation of Si at 1660 C. Interfacial vapor pressures have also been calculated for the Si compounds of interest. For SiC, interfacial vapor pressures have been calculated assuming the bounding cases of unit Si and unit C activity. These activity values result in 1.01 × 105 Pa interfacial pressure (sum of all vapor species) at temperatures of approximately 1800 and 1515 C, respectively. Calculations for Si3N4 yield a 1.01 × 105 Pa interfacial pressure at 1790 C for the case of unit Si activity. For MoSi2 , 1.01 × 105 Pa interfacial vapor pressure is predicted to occur at 1850 C for the case of unit Si activity [73]. Experimental data for these three SiO2 -forming compounds have shown reasonable agreement with predictions [74-77]. It is not clear that the equilibrium interfacial vapor pressure for any pure SiO2 -forming compound can be low enough to significantly increase the use temperature above approximately 1800 C (Jacobsen [65] has set the temperature limit at approximately 1725 C). Improved performance appears to be possible only by reducing the Si activity. This path, however, is limited since a lower Si activity will introduce additional components to the oxide scale and compromise the protective capability of pure SiO2 . It is emphasized here that this approximate 1800 C temperature limit is based on thermochemical quantities, and not on oxidation kinetics. It is conceivable that the other slow-growing oxides (based on Al, Cr, and Be) could have interfacial vapor pressures that are lower than for Si at high temperatures. Volatility diagrams for the Al-Al2O3 , Cr-Cr2O3 , and Be-BeO systems at 2227 C are shown in Figs 4-6, respectively (“JANAF” thermodynamic data [48] were used for all three metals). The diagrams reveal that the highest vapor pressures for these systems also exist at the metal-oxide interface. The vapor pressures for Al, Cr, and Be reach 1.01 × 105 Pa at 2490, 2690, and 2495 C, respectively. The vapor pressures were also computed as a function of temperature and are summarized in Table I. However, oxidation studies for materials based on these metals or compounds have revealed lower tem             Figure 4 Volatility diagram for Al-Al2O3 system at 2227     C.  5892  Figure 5 Volatility diagram for Cr-Cr2O3 system at 2227     C.  Figure 6 Volatility diagram for Be-BeO system at 2227     C.        perature limits than those imposed by the 1.01 × 105 Pa interfacial pressure. In isothermal oxidation of Cr at 980 C, an oxide scale was grown separated from the metal surface [78]. The scale grew by metal vapor transport from the oxide-free metal surface to the interior of the detached scale. Both Cr and Al have exhibited similar oxide scale disruption resulting from the high metal vapor pressures beneath the scale [79]. Oxidation data for beryllides show that protective scales have not been observed above 1650 C [80]. An interfacial vapor −4 Pa has been reported to be sufficient to pressure of 10 disrupt protective oxide scale formation [68, 77]. Table I also provides the calculated temperatures for which −4 Pa. It is noted that the interfacial vapor pressure is 10 these temperatures, 750 C for Si up to 970 C for Cr, are very low relative to the desired 2000 C capability. a 1.01 × 105 Pa interfacial vapor pressure (1800 The SiO2 -forming materials may be operated up to C), yet Cr2O3 -formers show unusual oxidation scaling be−4 Pa (at havior even at the interfacial pressure of 10 980 C). This may be explained, in part, by the differences in the mass transport mechanisms of the condensed oxides. The growth of SiO2 -scales has been correlated to the transport rate of molecular oxygen (O2 ) [65, 66, 80, 81]. The open structure of SiO2 glass also allows oxidation products (SiO, CO) at modest pressures to move outward without failure of the protective scale. Only when the interfacial pressure of these product species approaches 1 atm (or a lower pressure ambient environment for hypersonic applications), the                 \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  4. Theoretical aspects of the oxidation of boride materials     The UHTC materials based on mixtures of the ZrB2 or HfB2 and SiC have exhibited relatively good oxidation behavior in arc-heater testing up to the melting points of the base materials. Additionally, the oxidation data shown in Fig. 7 [82] do not reveal a catastrophic increase in oxidation rate from increasing interfacial vapor pressures. Fig. 8 shows a Scanning Electron Microscopy sis of the oxide scale cross-section on a ZrB2 + 20 vol% (SEM)/Energy Dispersive Spectroscopy (EDS) analySiC ceramic sample after furnace oxidation at 1600 C in air for 2 h. The backscatter image reveals the two phases (ZrB2 , SiC) in the unoxidized zone at the bottom of the photo. The EDS analysis reveals the presence of Zr and Si through the base material (the distribution of B and C were not conclusively identified by this EDS analysis). In the oxidized regions (four upper zones of the photo), Zr is the dominant element at the base of the scale, while the Si is the dominant element at the scale surface. Volatility diagrams provide insight into the oxidathese ZrB2 + 20 vol% SiC materition response of als. The diagram for the B-B2O3 system at 2227 C is shown in Fig. 9. The system is unique for the lack of a peak pressure at the B-B2O3 interface, and for the constant pressure through the oxide scale, compared to the Si-SiO2 , Al-Al2O3 , Cr-Cr2O3 , and Be-BeO systems (Figs 3-6). The volatility diagram for the Zr-ZrO2 system at 2227 C is shown in Fig. 10. The diagram is very similar     Figure 7 Arc-heater and furnace oxidation scale thickness versus tem perature results for ZrB2 -SiC.  system experiences disruptive degradation of the protective scale. For Al, Cr, and Be, the oxide scale is crystalline, which generally allows only ionic transport. For these metals, oxide scale growth includes, as a significant diffusion mechanism, metal cations diffusing outward from the metal-oxide interface to the oxide scale outer surface. Since molecular vapor species do not diffuse easily through a compact oxide scale, the scale is disrupted at relatively low vapor pressures. By this reasoning, materials that form a glass component in the scale are better choices for high temperature environments since they are more structurally tolerant of high interfacial vapor pressures.     Figure 8 Microstructure and elemental distribution in ZrB2 -SiC ceramics furnace-oxidized at 1600     C for 2 h.  5893  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 9 Volatility diagram for B-B2O3 system at 2227     C.  Figure 10 Volatility diagram for Zr-ZrO2 system at 2227     C.  to that for Si-SiO2 (Fig. 3), but exhibits significantly lower vapor pressures for all species. The maximum vapor pressure is found at the metal-metal oxide interface, but is of the order of 10 Pa compared to approximately 106 Pa for the Si-SiO2 interface. A combined volatility diagram of the Zr-ZrO2 , SiSiO2 , and B-B2O3 systems at 2227 C is shown in Fig. 11. The vapor pressure magnitudes will differ for the ZrB2 + 20 vol%SiC ceramics since unit activities are assumed for the calculation shown by Fig. 11. The diagram reveals significant insights into the be    Figure 11 Combined volatility diagram for Zr-ZrO2 , Si-SiO2 , and BB2O3 systems at 2227 C.     5894           havior of these materials. The vapor pressures associated with Zr (and Hf) are low enough to sustain an adherent oxide scale, which is confirmed by arc-heater testing. The pressures associated with B2O3 , although high, will be significantly lower due to the reduced B activity in the ZrB2 or HfB2 [83]. In addition, in a temperature gradient environment, such as in an archeater test (815 C gradients have been measured across a 2.5 mm scale thickness [82]) or on a hypersonic leading edge, the B2O3 vapor pressure will continuously decrease from the scale surface inward. Boria evaporation will then occur from the outer surface and will not catastrophically disrupt the oxide scale. Due to these temperature and pressure gradients, oxidation rates in arc heater testing are significantly lower, by up to 90% at 2200 C, than those measured using RF heating techniques [16, 19, 21]. This significant difference is also expected between arc-heater testing and furnace testing. The substantial temperature and Bx O y vapor pressure gradients through the scale promotes the retention of liquid B2O3 in the scale even to very high surface temperatures. Additionally, the B2O3 (l) wets the ZrO2 scale and persists due to the high surface energy of ZrO2 . At 1400 C, 10% of the B2O3 (l) formed is retained in the scale on pure ZrB2 [84]. For the ZrB2 -SiC ceramics, this retained B2O3 content is much higher due to formation of borosilicate glass. Gaseous SiO forms at the scale interface from SiC and migrates outward. Since O2 pressure also increases outward, SiO re-oxidizes to form condensed SiO2 at the exterior of the scale where it combines with B2O3 to form borosilicate glass. Since the B2O3 preferentially evaporates from the scale surface, the outer glass becomes enriched with SiO2 . The ZrO2 skeleton provides a framework for the glass to be retained and not removed by shear forces. The application use temperature of this materials system appears to be limited by melting of the oxide scale and/or the base material than by disruptively high interfacial vapor pressures [19, 21, 82]. Thus, these ZrB2 -SiC (and HfB2 -SiC) UHTC materials provide relatively good oxidation resistance by forming thermodynamically compatible oxide scale components. This scale system mitigates the effects of, or recovers from, high interfacial vapor pressures. However, the protection gained by the formation of the exterior SiO2 layer also provides the condition for P SiO to increase and become disruptive. As the SiO2 rich glass increases in thickness, O2 transport decreases, O2 pressure beneath the glass layer decreases, and PSiO increases until the glass layer is ruptured. New formation of SiO2 from SiO vapor suggests a cyclic protective/non-protective scale-forming sequence. A semi-protective scale should result, and this mechanism change in oxidation kinetics should be observed above the SiC-SiO2 active-passive transition temperature. It is not clear that this cyclic, semi-protective oxidation behavior has been observed to-date. The correlation between oxidation kinetics and oxygen transport mechanisms is not fully understood at this time. Up to approximately 1200 C, the parabolic rate constants for the oxidation of ZrB2 vary linearly with     \\x0c', '   P o2 , which affirms that molecular O2 transport through the B2O3 glass is the rate limiting step. Above 1200 C, the rate constants exhibit no dependence on P o2 , which is expected if oxygen transport through ZrO2 is the rate limiting step [79]. The addition of SiC has the effect of increasing the temperature for which glass transport properties provide the rate limiting step. The role of oxygen transport skeleton in ZrB2 + SiC materials is not clear at through the crystalline ZrO2 this time. Compositional variations, which form glasses that are stable to higher temperatures, and have low O2 diffusion rates and low vapor pressures, could further improve the oxidation resistance of these materials. Glasses have been identified which are composed of refractory oxides (La2O3 , ZrO2 , ThO2 , Ta2O5 , TiO2 , WO3 , and B2O3 ) [85]. It is recommended that new UHTC materials which exploit these glass-forming compositions be investigated. Compositional variations, which reduce oxygen transport through the skeleton phase, may improve the oxidation resistance as well. It is known that additions of higher valence metals into the ZrO2 lattice will reduce oxygen vacancy concentration and diffusion [79]. The pyrochlore phases associated with ZrO2 and rare earth oxides are also known to exhibit significantly lower oxygen diffusion than ZrO2 , as shown in Figs 12 and 13 [86, 87]. The in-situ formation of pyrochlore phases in UHTCs, and subsequent oxidation behavior, has not been investigated. It is recommended that new UHTC materials with rare earth additions be investigated. Vapor pressures at the metal-metal oxide interface are compared in Table I for potential skeleton phase-formers including Zr, Hf, Sc, Y, Be, and Ta. Materials forming HfO2 and ZrO2 are the best materials from this interfacial vapor pressure criterion. The carbides and nitrides of Zr (and Hf) are inferior to the diborides in oxidation behavior. Arc-heater testing of the carbides of Zr and Hf has been conducted at temperatures of 2400 to 2700 C with exposure times of 30 to 180 s in a gas environment replicating stoichiometric     Figure 12 Oxygen diffusivity for ZrO2 -rare earth oxide ceramics.  ULTRA-HIGH TEMPERATURE CERAMICS  Figure 13 Oxygen permeability constants for oxide ceramics with py rochlore structure.           hydrocarbon combustion [88]. A polished cross-section of HfC oxidized in arc-heater testing at 2700 for 30 s is shown in Fig. 14, revealing the porous oxide scale [89]. Vapor pressures for CO formed at the ZrC-ZrO2 and HfC-HfO2 interfaces have been calculated and exceed 1.01 × 105 Pa at 1730 C for both systems [90]. Thus, the carbides of Zr and Hf are compromised by the high vapor pressures formed at the oxide-carbide interface. The formation of porous scales was similarly observed for the nitrides. The formation of Zr and Hf oxycarbide intermediate phases in the ZrC-ZrO2 and HfC-HfO2 systems have been reported, as the primary oxygen barriers [91, 92] and as stable phases in themselves [93, 94]. Hafnium carbide was reported as exhibiting superior oxidation resistance to HfB2 at 1400 to 2100 C [91], but is based on furnace testing which biases the results unfavorably against the borides, as discussed previously. Stable oxycarbide phases have been reported for some rare earth metals, which suggests that rare earth additions may be employed to stabilize the oxycarbide phase in the scale. It is suggested that scandium (or ScC1−x ) be investigated since Sc has the same atomic radius as Zr or Hf and should exhibit high solubility into the base materials ZrC or HfC. Carbon interlayers, forming between the oxide and the base material, have also been reported for these carbides [95, 96]. Finally, the carbides are vulnerable to forming a powdery and non-adherent oxide scale below approximately 1700 C, which has been attributed to an inability to sinter at these temperatures [29, 97, 98]. Sintered and zone-melted samples exhibited low-temperature     5895  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 14 Microstructure (polished cross-section) of HfC arc-heater oxidized at 2700  C for 30 s.     catastrophic oxidation [30], while chemically vapor deposited (CVD) materials have exhibited protective scale formation [91]. The nature of the low temperature oxidation vulnerability is not understood at this time.  5. Oxidation studies—Glass scale modification  Oxidation behavior of non-oxide ceramics depends highly on the properties of the oxidation product and on the combination of physical and chemical processes taking place on the surface exposed to oxygencontaining atmosphere. In general, the chemical composition and structure of an oxidized surface define the oxidation stability of a ceramic material. Modification of the chemical composition of the oxide surface layer, leading to decreased inward diffusion of oxygen, is one of the efficient ways of controlling oxidation resistance of non-oxide ceramics. This modification can be accomplished by changing the bulk composition, or the  5896  r     surface of ceramics using CVD, ion implantation, “pack cementation”, and other methods. For example, the oxidation resistance of ZrB2 ceramics was significantly improved by modifying the bulk composition with SiC leading to the formation of a protective surface layer of borosilicate glass during exposure to an oxygen-containing atmosphere [18, 20, 33, 34, 39, 84]. The research conducted at NSWCCD showed that a further improvement of the oxidation performance of ZrB2 -SiC ceramics (up to 1600 C to-date) can be accomplished by the addition of CrB2 , TiB2 , TaB2 , NbB2 , and VB2 [99]. Oxidation of the modifying diborides resulted in the formation of corresponding oxides in the surface borosilicate glass. It should be emphasized that the oxidation resistance of all of the modifying diborides (alone) is much lower than that of the ZrB2 and HfB2 ceramics [14, 22]. It is known from the literature that borate and silicate glasses containing Group IV-VI transition metal oxides show strong tendency to phase separation (immiscibility) [100]. Systems exhibiting immiscibility are characterized by steeply rising liquidus temperatures and increased viscosity. An increase in the viscosity decreases the oxygen diffusion rate through the oxide surface scale based on the Stokes—Einstein relationship [101], which shows that diffusivity is inversely proportional to viscosity. Another potential benefit of increased viscosity as well as increased liquidus temperature is the suppression of boria evaporation from the glass. The oxide effectiveness in enhancing immiscibility increases with increasing metallic element cation field strength, z /r 2 , where z is the valence and is the ionic radius [100-103]. Since the cation field strengths of Ti, Nb, Ta, Mo, Cr, and V are higher than that of Zr, these elements can be effective in promoting phase separation of the borosilicate glass formed on the surface of the ZrB2 -SiC ceramics. The concept of phase separation as a controlling factor in the oxidation protection of non-oxide ceramics is unique and has not been discussed in the literature. The modifying diboride additives were introduced into the ZrB2 -SiC ceramics in the amounts of 2- 20 mol% as a substitution for ZrB2 . The molar ratio of ZrB2 to SiC in all materials was maintained at 2 (25 vol% SiC). The ceramics were prepared by hot pressing starting mixtures, consisting of ZrB2 , SiC, and modifying additives, at 2100 C and 20 MPa for 0.5 h. Oxidation experiments were conducted by heating the samples in a furnace in air at 1200-1600 C, typically for 2 h. The samples were placed into the furnace at the test temperatures and then air quenched after the hold. Quenching of the samples was conducted to retain the high temperature condition of the surface layer for analysis. Additionally, the oxidation of the samples was characterized by thermal gravimetric analysis (TGA) during 5 h isothermal heating at different temperatures in the air-simulating oxygen/argon mixture. The composition and structure of the surface and crosssection of the oxidized ceramics were evaluated using X-ray diffraction (XRD), SEM, and EDS. Fig. 15 shows the results of isothermal (5 h) TGA heating at 1300 C of ZrB2 -SiC ceramics modified with           \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS     SEM studies showed the evidence of hightemperature phase separation on the surfaces of all the modified samples after both TGA and furnace oxidation tests. Fig. 16a shows the surface of the sample containing 10 mol% TaB2 after the TGA oxidation test. The presence of large (more than 100 µm) droplets of borosilicate glass periodically distributed in a partially crystallized glassy matrix is an indication of high temperature glass phase separation. The matrix glass is enriched with Zr and Ta (from the EDS analysis data). X-ray diffraction of the whole surface of the oxidized sample showed the presence of ZrO2 along with a small amount Ta Zr2.75O8 with melting temperature above 1700 C. The CrB2 -containing ceramics (Fig. 16b) exhibited multiple phase separation after furnace oxidation at 1500 C, with shells around chromium-rich droplets which is the consequence of incomplete diffusion during cooling. A periodic pattern of crystals on the surface of ceramics modified with NbB2 (Fig. 16c) implies the existence of glasses of different compositions at the test temperature. X-ray diffraction analysis of the surface showed that the crystalline phase is Nb2Zr6O17 with a melting temperature about 1500 C. The glassy phase contains only small amounts of Nb and Zr (from EDS analysis). The presence of glass immiscibility is also observed on the surface of the ceramics modified with 5 mol% VB2 after TGA oxidation at 1300 C (Fig. 16d). Circular microcracking around droplets in this sample           Figure 15  C of ZrB2 /SiC ceramics modified with 10 mol% CrB2 , TiB2 , NbB2 , VB2 , and TaB2 .  Isothermal TGA oxidation at 1300     10 mol% of CrB2 , NbB2 , TaB2 , TiB2 and VB2 . The presence of the corresponding oxides in the oxidized surface layer improved the oxidation performance of the baseline material. The lowest weight gain during oxidation was observed for the ceramics containing TaB2 . The thickness of the oxidized layer of this sample was less than half of that for the baseline material [99]. The weight gain decreased in the sequence of the modifying additives: CrB2 , TiB2 , NbB2 , VB2 , and TaB2 , correlat+3 , Ti +4 , ing well with the cation field strength for Cr +5 , V +4 , and Ta +5 [99]. Nb  Figure 16 SEM micrographs of the surface of oxidized ZrB2 -SiC ceramics modified with: (a) TaB2, (b) CrB2 , (c) NbB2 , and (d) VB2 .  5897  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 17  C of TiB2 and TiB2 -SiC ceramics modified with 20 mol% ZrB2 , CrB2 , NbB2 , and TaB2 .  Isothermal TGA oxidation at 1300     can be attributed to compositional differences resulting in thermal expansion coefficient mismatch and leading to the development of strain and cracking during cooling. X-ray diffraction showed the presence of V7O13 in addition to ZrO2 on the surface of the oxidized VB2 containing sample. The discovered correlation between oxidation resistance of ZrB2 and the presence of phase separation in the surface protective glass was successfully applied to improve the oxidation performance of CrB2 , TiB2 , TaB2 , NbB2 , ZrB2 -Si3N4 , Ti3SiC2, and Si3N4 ceramics. The major results for the TiB2 , Ti3SiC2 , and Si3N4 ceramics are discussed below. The oxidation behavior of the TiB2 ceramics modified with SiC and 5-20 mol% CrB2 , NbB2 , TaB2 , and ZrB2 was evaluated [104]. The ceramics were prepared by hot pressing starting mixtures consisting of TiB2 ,        SiC, and diboride modifiers at 2100 C and 20 MPa for 0.5 h. Both furnace heating and TGA were used to characterize the oxidation resistance of the ceramics. The data show that the TiB2 -SiC ceramics containing TaB2 , NbB2 , and CrB2 have the best oxidation performance at all temperatures and additive contents. Fig. 17 presents the results of isothermal TGA heating at 1300 C. The addition of ZrB2 led to the decrease in oxidation resistance, especially, at 20% loading. The effect of additives on the oxidation resistance of TiB2 ceramics correlates with the cation field strength values for the elements being highest +4 oxide, with cation field strength for Ta and Nb. The presence of Zr +4 , led to the low protecting capalower than that of Ti bilities of the surface glass. If the rule of mixtures, not cation field strength of elements, was a controlling factor in the properties of surface glass and its protective capabilities, the improvement in the oxidation behavior of TiB2 -SiC ceramics could be expected with the introduction of ZrB2 , having the highest oxidation resistance of all the studied diborides. The optical micrograph of the surface of the CrB2 -containing ceramics (Fig. 18) clearly shows the immiscibility of the glass with the Cr2O3 -rich green areas and TiO2 -rich brown areas. The Ti3SiC2 ceramics recently attracted considerable attention because of their unique microstructure resulting in the exceptional combination of properties such as a high melting point, high fracture toughness and thermal shock resistance, plasticity at elevated temperatures, high modulus, low hardness, easy machinability, and self-lubrication [105-107]. These properties make the Ti3SiC2 ceramics of very high practical importance in numerous applications. However, these ceramics experience significant oxidation at temperatures above 1000 C preventing their high-temperature application     Figure 18 Optical micrograph of the surface of TiB2 -SiC ceramics containing 20 mol% CrB2 after oxidation at 1200     C for 2 h.  5898  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  in oxidizing environments. The modification of the ceramics with 10 mol% of TaB2 , ZrB2 , CrB2 , NbB2 , and VB2 was undertaken to create phase-separated borosilicate glass on the surface of the ceramics and, thus, increase their oxidation resistance. The Ti3SiC2 ceramics were synthesized and densified by reactive hot-pressing at 1500 C and 20 MPa for 1 h in He using the novel displacement reaction:     Ti5Si3 + 4TiC + 2C → Ti3SiC2  The additives were introduced in the starting Ti5Si3 TiC-C mixtures. The samples were oxidized during air Figure 20 Oxidation of Ti3 SiC2 ceramics modified with 10 mol% TaB2 , ZrB2 , CrB2 , TiB2 , and NbB2 during furnace heating for 2 h.              furnace heating at 1000 to 1500 C, typically for two hours. The hot-pressed ceramics were fully dense and contained Ti3SiC2 , small amounts of TiC, and additional phases associated with the additives. The SEM micrographs (Fig. 19a) show the presence of high aspect ratio plate-like grains, which define the mechanical behavior of the ceramics [105-107]. The additives (TaB2 in Fig. 19b) did not noticeably change the microstructure of the baseline material. The unmodified ceramics started to oxidize substantially at 1100 C (Fig. 20). Of all the diboride additives, NbB2 and, especially, TaB2 significantly improved the oxidation behavior of the ceramics. After 2 h oxidation at 1400 C, the weight gain was 430 g/m2 for the baseline ceramics compared to 60 g/m2 for the ceramics containing 10 mol% TaB2 . After heating at 1300 C the thickness of the oxidized layer was about 250 and 50 µm for the baseline and 10 mol% TaB2—modified ceramics, respectively. The microstructure of the oxidized surfaces of the baseline and TaB2 -modified ceramics (Fig. 21) is notably different. A bimodal distribution of the TiO2 crystals with significantly different morphology and size is observed in the TaB2 -containing sample. The bimodal distribution and clustering of these crystals are probably an indication that crystallization occurred from immiscible glasses of different composition during hightemperature exposure. The pronounced effect of TaB2 and NbB2 forming corresponding oxides in the glass is related to the highest cation field strength of Nb and Ta compared to all other elements tested. Silicon nitride is one of the most promising candidates for high-temperature structural applications, such as hot section components for advanced gas turbines and high-efficiency microturbines. High-temperature applications of Si3N4 ceramics depend, to a very high degree, on their behavior in corrosive environments and, primarily, on their resistance to oxidation. The effect of transition-metal diborides (CrB2 , TaB2 , and ZrB2 ) was studied on the oxidation behavior and microstructure of oxidized Si3N4 ceramics containing 5 wt% Y2O3 and 2 wt% Al2O3 as sintering aids [108]. The ceramics were hot-pressed at 1825 C and 20 MPa  Figure 19 SEM micrographs of the fracture surface of Ti3 SiC2 ceramics: (a) baseline Ti3 SiC2 and (b) Ti3 SiC2 modified with 10 mol% TaB2 .     5899  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 21 SEM micrographs of the surface of Ti3 SiC2 ceramics after oxidation at 1200 with 10 mol% TaB2 .     C for 2 h: (a) baseline Ti3 SiC2 and (b) Ti3 SiC2 modified  5900  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS           in He for 1 h. The oxidation behavior was characterized after furnace heating at 1200-1600 C in air. Fig. 22 shows that no significant oxidation (weight change) was observed for any of the materials below 1300 C. However, above 1350 C, only CrB2 significantly increased the oxidation resistance of the Si3N4 ceramics. The effect of the additives increases with increasing temperature. The SEM observations showed that the oxidized surface of the CrB2 -modified ceramics is almost fully crystallized, while the surface of the baseline ceramics is covered by a poorly crystallized, phase-separated glass (Fig. 23). The XRD analysis identified α -cristobalite as the only crystalline phase on the surface of the baseline material. In contrast, yttrium disilicate (monoclinic β -Y2Si2O7 ) together with a small amount of α -cristobalite were found on the surface of the Cr-containing material. No Cr-containing  Figure 22 Oxidation of Si3N4 ceramics modified with 10 mol% ZrB2 , TaB2 , and CrB2 .  Figure 23 SEM micrographs of the surface of Si3N4 ceramics after oxidation at 1550 vol% CrB2 .     C for 2 h: (a) baseline Si3N4 and (b) Si3N4 modified with 5  5901  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS     compounds were detected, however, on the surface of the materials by either EDS or XRD. It is hypothesized that chromium oxide present in the glass during oxidation additionally contributes to phase separation in the glass with the formation of a B-O-Si droplet phase and a Al-Y-Cr-Si-B-O matrix phase. Homogeneous nucleation and the formation of Cr2O3 crystallization centers occur in the matrix phase, followed by the catalytic crystallization of elongated Y2Si2O7 grains, while a small amount of fine cristobalite grains tration of Y2O3 · 2SiO2 crystals (melting temperature crystallizes from the droplet phase. The high concen1775 C) on the surface of the Cr-containing ceramics provides effective oxidation protection during exposure to oxidizing atmosphere. The highest oxidation resistance was shown by the ceramics containing less than 5 vol% CrB2 . The results of this research clearly showed that the oxidation resistance of non-oxide ceramics can be significantly enhanced by compositional design leading to the formation of a surface layer of immiscible multicomponent glass. The resulting increased liquidus temperatures and viscosities, as well as decreased oxygen diffusivities, in the immiscible glasses are considered responsible for the observed improvement in the oxidation resistance of the ceramics. The difference in the immiscibility and the corresponding oxidation behavior for different ceramics and modifying additives is a function of the oxidation state of elements in a particular glass, which highly affects their cation field strength and, consequently, the tendency of the glass toward phase separation. For example, Cr exhibited different oxidation states depending on the base ceramic material, and test temperature and atmosphere (TGA versus furnace oxidation). The presence of CrO2 was detected in the oxidation layer of several samples. The effective cation field strength of +4 is much higher than Cr +3 (1322 compared to 793 Cr −2 ) while that of Ta +5 is constant at 1220 nm −2 . This nm offers an explanation for the alternating effectiveness of Ta and Cr modifiers in enhancing oxidation resistance.  6. Conclusions     High vapor pressures at the metal-metal oxide interfaces of the slow-growing oxides (SiO2 , Al2O3 , Cr2O3 , BeO) are disruptive at 1800 C or below. The high interfacial vapor pressures of these systems primarily result from the system thermochemistry and are only secondarily dependent on the oxidation kinetics. Among these oxides, a glass-forming SiO2 scale exhibits significantly higher temperature capability, compared to materials which form a crystalline Al2O3 , Cr2O3 , or BeO scale, due to the greater structural tolerance of glass to high interfacial vapor pressures. Materials that form a multi-component oxide scale, composed of a refractory oxide skeleton and an amorphous (glass) oxide component, provide good oxidation performance at hypersonic use temperatures up to, and above, 2000 C. This multi-component oxide system is the only structure known at this time that mitigates, or recovers from, high interfacial vapor pressures. The     5902     oxidation resistance of ZrB2 -SiC and other non-oxide materials is improved, to at least 1600 C, by compositional modifications with transition metal additives that promote immiscibility in the glass component of the scale. The oxidation mechanisms of materials forming this scale structure (e.g., ZrB2 -SiC) are still not wellunderstood despite 40 years of research.  Acknowledgements  Sustained support by Dr. Stephen Fishman of the Office of Naval Research, and by the Naval Surface Warfare Center (Dahlgren and Carderock Divisions), is gratefully acknowledged. The authors would like also to thank Dr. E. Opila for insightful editing contributions, Dr. B. Varshal for valuable discussions on glass immiscibility issues, Dr. D. Dallek for performing TGA tests, and intern students for participation in the experiments. One of the authors (M.M.O.) expresses deep appreciation for the privilege to learn metallurgical thermodynamics from Professor Robert Rapp of Ohio State University. The errors are the author’s, however.  References  1. M . A . Conference on Refractory Metals and Alloys, Chicago, April 1962,  L E V I N S T E I N , in Proceedings of Metallurgical Society  edited by M. Semchysen and I. Perlmutter (Interscience Publishers,  New York, London, 1963) p. 269.  2.  “High-Temperature  Inorganic Coatings,”  edited by J. Huminik  (Reinhold, New York, 1963).  3.  “Coatings of High-Temperature Materials,” edited by H. H. Haus ner (Plenum, New York, 1966).  4.  “Protective Coatings on Metals,” edited by G. V. 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Chem. 19 (1993) 1.  5903  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  104.  J .  A .  Z A Y K O S K I  et al., FY 2000 NSWC Res. Digest  (2000)  108.  I .  G .  T A L M Y et al.,  in Proceedings of the International Sym posium on “High Temperature Corrosion and Materials Chem istry IV,” edited by E. Opila, P. Hou, T. Maruyama, D. Shifler and  E. Wuchina (The Electrochemical Society, Inc., Pennington, NJ,  95.  105. M . W . B A R S O U M and T . E L R A G H Y , J. Amer. Ceram. Soc.  79 (1996) 1953.  J .  F R Y T and L .  106.  S T O B I E R S K I , Trans Tech Pubs., Switzerland,  2003) Vol. 2003-16, p. 361.  (1997) p. 1608. 107. N . (1999) 4385.  F . G A O , M I Y A M O T O and D .  Z H A N G , J. Mater. Sci. 18  Received 28 February and accepted 29 April 2004  5904  \\x0c', '\\x0c']"
},{
  "_id": 199,
  "PDF": "Oxidation-based materials selection for 2000◦C +hypersonic aerosurfaces.pdf",
  "Text": "['ULTRA-HIGH TEMPERATURE CERAMICS  J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5887 - 5904  Oxidation-based materials selection for 2000 C + hypersonic aerosurfaces: Theoretical considerations and historical experience  M . M . O P E K A , I . G . T A LM Y , J . A . Z A Y K O S K I Naval Surface Warfare Center Carderock Division, West Bethesda, MD 20817, USA E-mail: opekaMM@nsweed.navy.mil  Hypersonic flight involves extremely high velocities and gas temperatures with the attendant necessity for thermal protection systems (TPS). New high temperature materials are needed for these TPS systems. A systematic investigation of the thermodynamics of potential materials revealed that low oxidation rate materials, which form pure scales of SiO2 , Al2O3 , Cr2O3 , or BeO, cannot be utilized at temperatures of 1800 C (and above) due to disruptively high vapor pressures which arise at the interface of the base material and the scale. Vapor pressure considerations provide significant insight into the relatively good oxidation resistance of ZrB2 and HfB2 -based materials at 2000 C and above. These materials form multi-oxide scales composed of a refractory crystalline oxide (skeleton) and a glass component, and this compositional approach is recommended for further development. The oxidation resistance of ZrB2 -SiC and other non-oxide materials is improved, to at least 1600 C, by compositional modifications which promote immiscibility in the glass component of the scale. Other candidate materials forming high temperature oxides, such as rare earth compounds, are largely unexplored for high temperature applications and may be attractive candidates for hypersonic TPS materials. C(cid:2) 2004 Kluwer Academic Publishers  1.  Introduction     The 21st century has ushered in a new, exciting era of hypersonic flight. Hypersonic flight vehicles include sub-orbital and earth-to-orbit vehicles for rapid global and space access missions. A common aspect of these future systems is the need for new high-temperature materials. Hypersonic vehicles with sharp aerosurfaces, such as engine cowl inlets, wing leading edges (LEs), and nosecaps, have projected needs for 2000 to 2400 C materials which must operate in air and be re-usable. At this time, there are few, if any, off-the-shelf materials to meet these future hypersonic thermal protection system (TPS) needs. State-of-the-Art high temperature materials include carbon-carbon composites (CC) and silicon carbide-based (SiC) composites, such as C-SiC and SiC-SiC. Ultra-High-Temperature Ceramics (UHTCs), such as Zr(Hf)B2 -SiC, are being developed but are less mature at this time. Carbon-carbon composites have very high temperature structural capabilities but are not oxidation-resistant. Coatings have been and are being developed for oxidation-resistance, but cyclic life capabilities are modest due to the difficulties of managing the thermal expansion coefficient (CTE) mismatch between the C-C composite and the coating systems. The SiC-based composites exhibit oxidation resistance up to 1600 C in hypersonic environments, but thermal cycling lifetimes are also modest due to CTEmismatch-induced matrix cracking which        allows direct oxidation of the carbon fiber reinforcement. The UHTCs, based on the diborides of zirconium and hafnium, have exhibited relatively good oxidation resistance above 1600 C. The oxidation mechanisms of these materials are not well understood. Recent Navy efforts to understand UHTC oxidation mechanisms, and to develop new, highly oxidation-resistant 2000 C materials, are presented here. This paper describes prior development of ultra-high-temperature, oxidation-resistant materials; thermodynamics and kinetics principles related to oxidation; theoretical aspects of the oxidation of UHTC materials; and experimental results associated with compositional variations of UHTC materials.     2. Developmental history  Distinct lines of research have contributed significantly to our current understanding of oxidationresistant ultra-high temperature materials: coating systems for refractory metals and subsequent development of oxidation-resistant intermetallic compounds, oxidation-resistant graphite compositions, and the development of boride-based UHTCs. The structural usefulness of refractory metals, and their lack of high temperature oxidation-resistance, motivated the pursuit of oxidation-resistant coatings. Considerable research was conducted, especially in the  0022-2461  C(cid:2) 2004 Kluwer Academic Publishers  5887  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS        1960s, and a number of texts summarizing the developments are available [1-5]. Although a broad range of materials was investigated, a significant proportion of the work was based on compositions containing silicon (Si), aluminum (Al), and chromium (Cr). Packer [6] summarized research on silicides, and an important conclusion is the significance of low pressure and high temperature environments on limiting the life of such materials and coatings. Perkins and Packer [7] identified the maximum temperature capability of MoSi2 coatings as 1800 C in atmospheric pressure (hypersonic) environments. Recent research on the oxidation of intermetallics, especially emphasizing aluminides for gas turbine applications, has been compiled by Grobstein and Doychak [8]. Oxidation-resistant graphite compositions were developed in parallel with refractory metal coatings in the 1960s [9, 10]. One of the most important compositions, designated “grade JTA” graphite [9], was optimized for oxidation resistance at 2000 C. It used additions of ZrB2 and Si (balance carbon) at an approximately 50 weight percent basis. Further optimization using transition metal additions (e.g., niobium) were found to improve oxidation performance at high temperatures, but with the penalty of poorer performance at lower temperatures. Krivoshein and coworkers [11] reported that Nb additions (10 wt%) improved oxidation performance of ZrB2 -SiC modified graphite, but that V additions at the same level provided maximum improvement. Significant research was reported on the refractory boride compounds beginning in the late 1940s with crystal structure [12] and melting point [13] measurements. An initial survey of the oxidation resistance of transition metal diborides up to 1500 C revealed that the group IVb compounds were the most resistant [14]. A survey of oxidation resistance of the diborides of Hf, Zr, Ti, Ta, and Nb from 1200 to 2200 C (inductively heated samples in flowing He-O2 mixtures) also revealed that HfB2 was the most oxidation resistant, followed by ZrB2 . The temperature dependence of the oxidation data for both compounds indicated significant rate changes at the respective metal oxide phase transition temperatures [15]. Oxidation testing in flowing He-O2 mixtures with H2O (at 613 Pa) exhibited a five-fold increase in oxidation rate of HfB2 at 933 C versus the nominal dry He-O2 . Similar oxidation measurements at 1487 C showed no rate difference [16]. Additional oxidation studies on ZrB2 and HfB2 demonstrated that metal-rich compositions (e.g., HfB1.7 ) oxidized at lower rates (by up to a factor of 50) versus boron-rich compositions (e.g., HfB2.12 ) [16]. Numerous investigations to improve the oxidation resistance of ZrB2 and HfB2 have been reported [17-20]. Compositions with 5 to 50 vol% SiC were investigated for both ZrB2 and HfB2 over a wide range of test temperatures and pressures; 20 vol% compositions were judged optimal for hypersonic vehicles in a series of efforts supported by the US Air Force [16, 19, 21-23]. Additions of C (5, 10, 15, 20, 30, and 50 vol%) improved thermal stress resistance, but were detrimental to oxidation resistance at all proportions. Additions of              5888     Cr (10 mol%), Al (20 mol%), and Ta (30 mol%) were found to be detrimental to oxidation resistance. An addition of 4 vol% of a Hf-20 at.%Ta alloy had no effect on the oxidation properties, although the metal phase was converted to the carbides during the hot-press fab50 mol% HfB2 + 50 mol% HfB composition exhibited rication process. Excess Hf metal additions to produce a rapid, preferential oxidation of the HfB phase. Additions of silicon to substitute on the boron sub-lattice yielded a HfB2 + “HfSi” skeletal phase which also exhibited rapid, preferential oxidation. (It is noted here that “HfSi” was identified by X-ray diffraction; other Hf-Si or Zr-Si second phases are possible, but have not been explored for oxidation response.) Additions of SiB6 (10 and 20 vol%) were found to increase oxidation resistance, but were not superior to SiC additions. Other systematic studies of additions into ZrB2 and or HfB2 have been conducted. Shaffer [24] evaluated the oxidation resistance of ZrB2 with additions of the disilicides of Ta, Nb, W, Mo, Zr, Mo0.5Ta0.5 , and Mo0.8Ta0.2 , as well as Zr5Si3 . The additive amounts were not specified, however, and only the conclusion was stated that MoSi2 was “unquestionably the best.” Additional oxidation experiments with varying proportions of MoSi2 (1 to 20 mol%) were conducted at 1950 C and revealed + 10 mol%MoSi2 composition was marketed by the the optimum composition to be 10 mol%. The ZrB2 Carborundum Company (US) under the Trade name “Boride Z”. Pastor and Meyer [25] evaluated the oxidation resistance of ZrB2 with additions of MSi2 or M5Si3 , where M is a transition metal Zr, Ta, Cr, Mo, or W. On the basis of scale thickness measurements after oxidation testing for up to 100 h at 1200 and 1400 + 15 wt%CrSi2 composition was found to be the most C, the ZrB2 oxidation resistant. Lavrenko and coworkers [26] reported that a ZrB2 + 50 wt%ZrSi2 composition was more oxidation resistant than MoSi2 and WSi2 , and could be used up to 1700 C. However, since oxidation data only up to 1200 C are reported, it is not clear how the conclusion is supported. The oxidation kinetics mechanism(s) of the diboridebased materials are only partly understood despite significant research. Oxidation kinetics measurements are typically based on weight change or scale thickness changes with time upon exposure to a known temperature and oxidizing atmosphere. However, weight change and scale thickness measurements are confounded by simultaneous oxidation and vaporization (of BOx vapor species) processes. Total oxygen consumption measurements (per unit area of sample) have been utilized to overcome this limitation [27]. Initial oxidation studies were conducted in 1955 on porous ZrB2 samples from 649 to 1315 C [28]. The oxidation kinetics were found to be parabolic, the rates increased with oxygen partial pressure, and the presence of H2O also increased the oxidation rate. Berkowitz-Mattuck measured total oxygen consumption for ZrB2 over a higher temperature range (1200-1530 C) and a lower oxygen partial pressure ( P o2 ) range (1070 to 5200 Pa) in helium (He) at 1.01 × 105 Pa total pressure [27]. Parabolic rate kinetics were                 \\x0c', '                  also observed, as were modest increases in oxidation rates with increasing P o2 . From metallographic examination of tested samples, it was concluded that oxidation proceeded by inward diffusion of oxygen, and it was suggested that oxygen diffusion through ZrO2 was the rate controlling step. Kuriakose and Margrave measured weight changes for ZrB2 over the temperature range of 945-1256 C and also reported parabolic oxidation kinetics [27, 29]. At 1056 C they observed that the parabolic rate con1.36 × 104 to 9.92 × 105 Pa at 1.01 × 105 Pa total presstant increased directly proportional to P o2 (range was sure with balance He). Berkowitz-Mattuck extended the oxidation kinetics studies of ZrB2 to understand the change in P o2 dependence with temperature [30, 31]. The P o2 dependence was confirmed at a test temperature of 927 C, but no dependence was found at 1557 C. Additional testing also revealed a significant change in the activation energy at 1057 C, changing from 20 kcal/mole below this temperature to 70 kcal/mole above it. Abrupt changes in the oxidation rate kinetics were also observed at the temperatures corresponding to the monoclinic to tetragonal oxide phase transitions for both ZrB2 and HfB2 . Other oxidation studies have been conducted on the oxidation kinetics of ZrB2 , HfB2 , and their respective SiC-modified compositions [32-40]. Oxygen diffusion through the B2O3 liquid phase was identified as the rate limiting step associated with oxidation of the pure diborides up to approximately 1200 C. Above this temperature, the increased oxidation rates were attributed to oxygen transport through the ZrO2 or HfO2 phase. The addition of SiC was found to significantly increase the temperature range of the glass as the primary oxygen barrier. A two layer scale was observed to form with HfO2 inner and SiO2 outer components. The reduction in oxidation rate was observed above 1350 C. Below this temperature, SiC inclusions are found in the HfO2 scale since the SiC particles do not oxidize significantly to generate the SiO2 glass component [34]. Low temperature oxidation studies have also been conducted for these borides. Preferential oxidation of C at an oxygen pressure of 1.3 × the Zr or Hf at 500 −3 Pa has been reported. Boron inclusions, which co10 alesced into layers, were observed in the oxide scale. Solution of oxygen into the diboride lattice was also reported [36]. Changes in the oxidation mechanism were noted at approximately 500 C [37]. In addition to the diborides, other materials were investigated for potential hypersonic applications [20, 41-44]. Materials based on ZrC and HfC were extensively studied, but were found to oxidize (nonprotectively) below 1800 C, which eliminated them from consideration for the temperature cycling hypersonic applications. Additions of SiC did not solve the rapid oxidation at low-temperatures. Hafniumtantalum alloys (e.g., Hf-20 at.%Ta) were found to exhibit good oxidation behavior, but were limited by the relatively low melting point of 2000 C at this composition. Iridium coatings on graphite were also evaluated, but were judged costly and not sufficiently refractory due to the Ir-C eutectic at 2296 C. The viscosity of                    ULTRA-HIGH TEMPERATURE CERAMICS     SiO2 was significantly increased by the addition of W powder in 10 and 20 vol% additions. These materials exhibited increased sensitivity to thermal stress failure and deformed by viscous flow into blunt shapes. Since hypersonic applications typically require the high temperature materials to retain sharp radii for leading edges and propulsion inlets, this shape change was unacceptable. By the early-1970s, the ZrB2 and HfB2 -based materials were identified to be the most promising for hypersonic applications with cyclic exposure from ambient temperature up to 2700 C [18, 41, 42]. Prior materials development for hypersonic applications does not include significant emphasis on oxide materials. They have not been pursued for these applications due to the demanding structural and thermostructural requirements of such systems, and low thermal shock resistance of oxides in general. It must be asked whether the optimized choice of hypersonic materials should be oxide or non-oxide materials. Oxide materials are, at best, intrinsically resistant to oxidation. However, oxide-oxide composites for 2000 C usage do not currently exist, and current and developmental oxide composites for aircraft applications cannot be used to that temperature. Such ultra-high-temperature oxide composites will likely be very costly due to the need to develop new creep-resistant reinforcements and suitable fiber-matrix interface materials to evade brittle fracture behavior. In addition, oxides and oxide composites incur significant design penalties due to their relatively high CTE and stiffness, and low thermal conductivity. Such a new 2000 C oxide composite would have to be developed for dedicated hypersonic application at very high cost. However, the payoff of intrinsic oxidation resistance requires a continuing look into oxide materials for these applications. The materials selection process presented here addresses the optimization of non-oxide ceramic compositions for high temperature hypersonic applications.        3. Oxidation-thermodynamics and kinetics  The selection of new oxidation-resistant materials is based on chemical thermodynamics and kinetics. Chemical thermodynamics is a powerful tool for identifying the equilibrium phases associated with the oxidation kinetics process(es). Chemical thermodynamics can be seen as providing the boundary conditions for understanding the oxidation kinetics processes. The thermodynamics-based calculations provided in this paper are based on the following relation:  \\x01G = \\x01G    + R · T · ln(Q )  (1)  where \\x01G is the change in Gibbs Free Energy (superscript refers to standard state) associated with a given chemical reaction, R is the ideal gas constant, T is the reaction temperature, and Q is the activity quotient [45, 46]. At the condition of chemical equilibrium, \\x01G is zero and Equation 1 reduces to:    = − R · T · ln(k )  \\x01G  (2)  5889  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  or,    = −2.303 · R · T · log(k )  \\x01G  (3)  where k is the reaction equilibrium constant. Gibbs Free Energies of Formation (FEOF) at standard state (\\x01G o f ) have been measured, and/or estimated, and tabulated as a function of temperature for many compounds of interest [47-50].  3.1. Thermodynamics and condensed phase equilibria  A stability diagram (Fig. 1) showing regions of metal and condensed metal oxide equilibria using Equation 3 was constructed for the elements Be, Y, Hf, Ta, W, Re, and Ir. These elements were selected as the most refractory representatives of their respective columns or groups in the periodic chart of elements. The diagram provides the metal-metal oxide equilibria as functions of temperature and oxygen pressure, P o2 . Each metal is stable below its respective metal-metal oxide equilibrium line, while the condensed oxide is stable above the equilibrium line. This diagram allows the generalized hypersonic environment to be directly compared with the metal-metal oxide stability regions. The FEOF data for BeO, HfO2 , Ta2O5 , and Re2O7 were taken from Schick [47], for WO3 from “JANAF” [48] , for Y2O3 from Pankratz [49], and for IrO2 from Knacke [50]. As necessary, FEOF were extrapolated linearly from the highest temperature data available. For Y2O3 , WO3 , Re2O7 , and IrO2 , the data were extrapolated above 1723, 2723, 361, and 1027 C, respectively, to compare with the hypersonic vehicle environment. Multiple condensed oxides exist and were considered for W (WO2 and WO3 ) and Re (ReO2 , ReO3 , and Re2O7 ). For both metals, the suboxides (WO2 , ReO2 , and ReO3 ) decompose below 1500 C and were neglected. The hypersonic environment envelope divides the diagram of Fig. 1 into three distinct groups of elements. The first group includes noble metals, such as iridium (Ir), for which the metal is the equilibrium condensed phase in the hypersonic environment. Since a condensed oxide does not form, mass loss from Irx O y        Figure 1 Metal-metal oxide condensed phase equilibrium diagram.  5890     vapor species must be considered. Oxidation mass loss rates for Ir are extremely low [51-53]. Other possible noble elements include rhodium, platinum, palladium, osmium, ruthenium, gold, and silver, which are not shown. Osmium and Ru exhibit very high mass loss rates in oxidizing environments [53]. Of the remaining elements, only iridium, and possibly rhodium, have melting points high enough to be considered for hypersonics applications. They should be considered for development but are not discussed further in this paper. The second element group shown in Fig. 1 includes those elements which form condensed oxides in the hypersonic environment, but the oxide melting point is below the 2000 C upper use temperature. These elements include tantalum, tungsten, rhenium, and by analogy, aluminum, titanium, molybdenum and all other group Vb (vanadium, niobium) and VIIb (manganese, technetium) elements. Since the oxides of these elements have relatively low melting points and melt viscosities, aerodynamic shear forces would remove them quickly and high ablation rates would be observed. In addition to the low oxide melting point limitation, these metals oxidize and experience high mass loss rates (from oxidative vaporization) at modest to high temperatures. Oxidation rates for Ta and W (as well as for Nb and Mo) have been reviewed by Kofstad [54]. High oxidation rates for Re in air are reported in the 300 to 1500 C temperature range [55, 56]. The third group, hafnium (Hf), yttrium (Y), and beryllium (Be), includes those elements which exist only as solid, condensed oxides throughout the hypersonic environment. Although only these three elements are shown, the result may be generalized to include zirconium, chromium, silicon, scandium, and the rare earth metals. Silicon is also included in this group because of the high viscosity of the molten oxide. Since these elements have melting points near to or lower than 2000 C, they have to be used as refractory compounds (silicides, borides, carbides, nitrides, etc.). Thus, initial candidate hypersonic materials include noble, refractory metals (Ir and/or Rh), and compounds (based on Zr, Hf, Be, Si, Y, Sc, and rare earths) which form refractory oxides.        3.2. Oxidation kinetics  While equilibrium thermodynamics provides an important foundation for materials selection, a second foundation is oxidation kinetics. For hypersonic applications, materials with a low oxidation rate are required. The need for higher temperature gas turbine materials has motivated significant investigations into very low oxidation rate materials. As shown in Fig. 2, Zr and Hf exhibit relatively high oxidation rates, while the most slowly oxidizing materials at high temperatures are those which form scales composed of pure SiO2 , Al2O3 , Cr2O3 , or BeO [57-64]. Above approximately 1100 C, SiO2 scale-forming materials (SiC, Si3N4 , MoSi2 , and possibly other silicides) have the lowest known oxidation rates [65, 66]. Beryllides (e.g., Ta2Be17 and ZrBe13 ) have very low oxidation rates up to 1250 C, but the rates increase rapidly with temperature [62, 63].        \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 2 Parabolic oxidation rate constants for various metals and com pounds.  3.3. Thermodynamics and vapor phase equilibria  Equation 3 may also be employed to construct diagrams of the vapor pressures of metal and metal oxide gaseous species as a function of oxidant pressure. Such a vapor phase diagram [67, 68], or volatility diagram, for the silicon-oxygen (Si-O) system at 2227 C is shown in Fig. 3. The vertical dashed line, at an oxygen pressure ( P o2 ) of 3.0×10 −5 Pa, is the equilibrium line defined by the reaction:     Si(c, l) + O2 (g) → SiO2 (l)  (4)  and separates the diagram into the two regions in which condensed Si and SiO2 exist. Unit activity is assumed for both Si and SiO2 , which is typical for this type diagram. Other lines show the vapor pressures of the SiO gaseous species (Si, Si2 , Si3 , SiO, SiO2 ) as a function of P o2 . The diagram may be seen as a schematic of the vapor pressure changes from the external surface of the SiO2 scale (right side of diagram) on oxidized Si, to the interior of the Si (left side of the diagram). It should be noted that the vapor pressures are highest at the Si-SiO2 interface (SiO vapor) and not at the exterior surface of the SiO2 . This high interfacial vapor pressure limits the use temperature of SiO2 -forming mate T A B L E I Calculated vapor pressures for oxides  Dominant  log P MOx  log PMOx  Condensed  Oxide  Melt temp  vapor species  at 1727     C  at 2227     C  for P = 10 Temp (deg C) −4 Pa  Temp (deg C) for P = 105 Pa  Material  oxide  type  (deg C)  (at Interface)  (Pa)  (Pa)  O2 Barrier matls Be  BeO  Crystalline  2550  Be  2.85  4.43  820  2495  Si  SiO2 Al2O3 Cr2O3 B2O3  Amorphous  1725  SiO  4.5  6.06  750  1865  Al  Crystalline  2040  Al, Al2O Cr  2.87  4.42  800  2490  Cr  Crystalline  2300  2.14  3.9  970  2690  B  Amorphous  450  B2O3 , B2O2  4.11  5.88  800  1950  Scale structures  Be  BeO  Crystalline  2550  Be  2.85  4.43  820  2495  Sc  Sc2O3 Y2O3 ZrO2 HfO2 Ta2O5  Crystalline  2400  Sc  2.02  3.8  950  2750  Y  Crystalline  2430  Y  −2.18 0.626 −2.58 −1.92  2.63  1120  3330  Zr  Crystalline  2700  ZrO  0.926  1520  3640  Hf  Crystalline  2800  HfO  0.506  1570  3670  Ta  Crystalline  1890  TaO2  0.926  1490  3730  Figure 3 Volatility diagram for Si-SiO2 system at 2227     C.  rials. The SiO2 scale is continuously ruptured when the interfacial (SiO) pressure exceeds ambient pressure and the scale loses its protective capability. At this condition, the SiO2 scale is consumed via the reaction:  Si(l) + SiO2 (l) → 2SiO(g)  (5)  The SiO(g) diffuses outward from the base of the porous SiO2 (l) scale and reacts with O2 (g) to form SiO2 (l), either within the scale or as smoke outside the scale [68]. If the SiO2 scale is entirely consumed (or if not formed initially), oxidation proceeds via the reaction:  Si(l) + 1/2O2 (g) → SiO(g)  (6)  The conditions represented by reactions (5) and (6) yield rapid oxidation of the Si. Wagner first showed that the transition between the rapid, (active) oxidation of Si(c,l) which occurs via reaction (6) and the slow (passive) oxidation via reaction (4) was governed primarily by the thermodynamics of the system [69]. Gulbransen includes reaction (5) in his description of active oxidation [68], although other researchers have restricted the term to describe the effect of reaction (6) only [70, 71]. The maximum in PSiO at the Si-SiO2 interface uniformly increases with temperature. The calculated SiO exceeds 1.01 × 105 Pa interfacial vapor pressure at 1865 C as shown in Table I.     5891  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS                    Gulbransen et al. [72] confirmed the active-passive transition boundary line for elemental Si in the 1100 to iments to air at 1.01 × 105 Pa yields the onset of active 1300 C temperature range. Extrapolating these experoxidation of Si at 1660 C. Interfacial vapor pressures have also been calculated for the Si compounds of interest. For SiC, interfacial vapor pressures have been calculated assuming the bounding cases of unit Si and unit C activity. These activity values result in 1.01 × 105 Pa interfacial pressure (sum of all vapor species) at temperatures of approximately 1800 and 1515 C, respectively. Calculations for Si3N4 yield a 1.01 × 105 Pa interfacial pressure at 1790 C for the case of unit Si activity. For MoSi2 , 1.01 × 105 Pa interfacial vapor pressure is predicted to occur at 1850 C for the case of unit Si activity [73]. Experimental data for these three SiO2 -forming compounds have shown reasonable agreement with predictions [74-77]. It is not clear that the equilibrium interfacial vapor pressure for any pure SiO2 -forming compound can be low enough to significantly increase the use temperature above approximately 1800 C (Jacobsen [65] has set the temperature limit at approximately 1725 C). Improved performance appears to be possible only by reducing the Si activity. This path, however, is limited since a lower Si activity will introduce additional components to the oxide scale and compromise the protective capability of pure SiO2 . It is emphasized here that this approximate 1800 C temperature limit is based on thermochemical quantities, and not on oxidation kinetics. It is conceivable that the other slow-growing oxides (based on Al, Cr, and Be) could have interfacial vapor pressures that are lower than for Si at high temperatures. Volatility diagrams for the Al-Al2O3 , Cr-Cr2O3 , and Be-BeO systems at 2227 C are shown in Figs 4-6, respectively (“JANAF” thermodynamic data [48] were used for all three metals). The diagrams reveal that the highest vapor pressures for these systems also exist at the metal-oxide interface. The vapor pressures for Al, Cr, and Be reach 1.01 × 105 Pa at 2490, 2690, and 2495 C, respectively. The vapor pressures were also computed as a function of temperature and are summarized in Table I. However, oxidation studies for materials based on these metals or compounds have revealed lower tem             Figure 4 Volatility diagram for Al-Al2O3 system at 2227     C.  5892  Figure 5 Volatility diagram for Cr-Cr2O3 system at 2227     C.  Figure 6 Volatility diagram for Be-BeO system at 2227     C.        perature limits than those imposed by the 1.01 × 105 Pa interfacial pressure. In isothermal oxidation of Cr at 980 C, an oxide scale was grown separated from the metal surface [78]. The scale grew by metal vapor transport from the oxide-free metal surface to the interior of the detached scale. Both Cr and Al have exhibited similar oxide scale disruption resulting from the high metal vapor pressures beneath the scale [79]. Oxidation data for beryllides show that protective scales have not been observed above 1650 C [80]. An interfacial vapor −4 Pa has been reported to be sufficient to pressure of 10 disrupt protective oxide scale formation [68, 77]. Table I also provides the calculated temperatures for which −4 Pa. It is noted that the interfacial vapor pressure is 10 these temperatures, 750 C for Si up to 970 C for Cr, are very low relative to the desired 2000 C capability. a 1.01 × 105 Pa interfacial vapor pressure (1800 The SiO2 -forming materials may be operated up to C), yet Cr2O3 -formers show unusual oxidation scaling be−4 Pa (at havior even at the interfacial pressure of 10 980 C). This may be explained, in part, by the differences in the mass transport mechanisms of the condensed oxides. The growth of SiO2 -scales has been correlated to the transport rate of molecular oxygen (O2 ) [65, 66, 80, 81]. The open structure of SiO2 glass also allows oxidation products (SiO, CO) at modest pressures to move outward without failure of the protective scale. Only when the interfacial pressure of these product species approaches 1 atm (or a lower pressure ambient environment for hypersonic applications), the                 \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  4. Theoretical aspects of the oxidation of boride materials     The UHTC materials based on mixtures of the ZrB2 or HfB2 and SiC have exhibited relatively good oxidation behavior in arc-heater testing up to the melting points of the base materials. Additionally, the oxidation data shown in Fig. 7 [82] do not reveal a catastrophic increase in oxidation rate from increasing interfacial vapor pressures. Fig. 8 shows a Scanning Electron Microscopy sis of the oxide scale cross-section on a ZrB2 + 20 vol% (SEM)/Energy Dispersive Spectroscopy (EDS) analySiC ceramic sample after furnace oxidation at 1600 C in air for 2 h. The backscatter image reveals the two phases (ZrB2 , SiC) in the unoxidized zone at the bottom of the photo. The EDS analysis reveals the presence of Zr and Si through the base material (the distribution of B and C were not conclusively identified by this EDS analysis). In the oxidized regions (four upper zones of the photo), Zr is the dominant element at the base of the scale, while the Si is the dominant element at the scale surface. Volatility diagrams provide insight into the oxidathese ZrB2 + 20 vol% SiC materition response of als. The diagram for the B-B2O3 system at 2227 C is shown in Fig. 9. The system is unique for the lack of a peak pressure at the B-B2O3 interface, and for the constant pressure through the oxide scale, compared to the Si-SiO2 , Al-Al2O3 , Cr-Cr2O3 , and Be-BeO systems (Figs 3-6). The volatility diagram for the Zr-ZrO2 system at 2227 C is shown in Fig. 10. The diagram is very similar     Figure 7 Arc-heater and furnace oxidation scale thickness versus tem perature results for ZrB2 -SiC.  system experiences disruptive degradation of the protective scale. For Al, Cr, and Be, the oxide scale is crystalline, which generally allows only ionic transport. For these metals, oxide scale growth includes, as a significant diffusion mechanism, metal cations diffusing outward from the metal-oxide interface to the oxide scale outer surface. Since molecular vapor species do not diffuse easily through a compact oxide scale, the scale is disrupted at relatively low vapor pressures. By this reasoning, materials that form a glass component in the scale are better choices for high temperature environments since they are more structurally tolerant of high interfacial vapor pressures.     Figure 8 Microstructure and elemental distribution in ZrB2 -SiC ceramics furnace-oxidized at 1600     C for 2 h.  5893  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 9 Volatility diagram for B-B2O3 system at 2227     C.  Figure 10 Volatility diagram for Zr-ZrO2 system at 2227     C.  to that for Si-SiO2 (Fig. 3), but exhibits significantly lower vapor pressures for all species. The maximum vapor pressure is found at the metal-metal oxide interface, but is of the order of 10 Pa compared to approximately 106 Pa for the Si-SiO2 interface. A combined volatility diagram of the Zr-ZrO2 , SiSiO2 , and B-B2O3 systems at 2227 C is shown in Fig. 11. The vapor pressure magnitudes will differ for the ZrB2 + 20 vol%SiC ceramics since unit activities are assumed for the calculation shown by Fig. 11. The diagram reveals significant insights into the be    Figure 11 Combined volatility diagram for Zr-ZrO2 , Si-SiO2 , and BB2O3 systems at 2227 C.     5894           havior of these materials. The vapor pressures associated with Zr (and Hf) are low enough to sustain an adherent oxide scale, which is confirmed by arc-heater testing. The pressures associated with B2O3 , although high, will be significantly lower due to the reduced B activity in the ZrB2 or HfB2 [83]. In addition, in a temperature gradient environment, such as in an archeater test (815 C gradients have been measured across a 2.5 mm scale thickness [82]) or on a hypersonic leading edge, the B2O3 vapor pressure will continuously decrease from the scale surface inward. Boria evaporation will then occur from the outer surface and will not catastrophically disrupt the oxide scale. Due to these temperature and pressure gradients, oxidation rates in arc heater testing are significantly lower, by up to 90% at 2200 C, than those measured using RF heating techniques [16, 19, 21]. This significant difference is also expected between arc-heater testing and furnace testing. The substantial temperature and Bx O y vapor pressure gradients through the scale promotes the retention of liquid B2O3 in the scale even to very high surface temperatures. Additionally, the B2O3 (l) wets the ZrO2 scale and persists due to the high surface energy of ZrO2 . At 1400 C, 10% of the B2O3 (l) formed is retained in the scale on pure ZrB2 [84]. For the ZrB2 -SiC ceramics, this retained B2O3 content is much higher due to formation of borosilicate glass. Gaseous SiO forms at the scale interface from SiC and migrates outward. Since O2 pressure also increases outward, SiO re-oxidizes to form condensed SiO2 at the exterior of the scale where it combines with B2O3 to form borosilicate glass. Since the B2O3 preferentially evaporates from the scale surface, the outer glass becomes enriched with SiO2 . The ZrO2 skeleton provides a framework for the glass to be retained and not removed by shear forces. The application use temperature of this materials system appears to be limited by melting of the oxide scale and/or the base material than by disruptively high interfacial vapor pressures [19, 21, 82]. Thus, these ZrB2 -SiC (and HfB2 -SiC) UHTC materials provide relatively good oxidation resistance by forming thermodynamically compatible oxide scale components. This scale system mitigates the effects of, or recovers from, high interfacial vapor pressures. However, the protection gained by the formation of the exterior SiO2 layer also provides the condition for P SiO to increase and become disruptive. As the SiO2 rich glass increases in thickness, O2 transport decreases, O2 pressure beneath the glass layer decreases, and PSiO increases until the glass layer is ruptured. New formation of SiO2 from SiO vapor suggests a cyclic protective/non-protective scale-forming sequence. A semi-protective scale should result, and this mechanism change in oxidation kinetics should be observed above the SiC-SiO2 active-passive transition temperature. It is not clear that this cyclic, semi-protective oxidation behavior has been observed to-date. The correlation between oxidation kinetics and oxygen transport mechanisms is not fully understood at this time. Up to approximately 1200 C, the parabolic rate constants for the oxidation of ZrB2 vary linearly with     \\x0c', '   P o2 , which affirms that molecular O2 transport through the B2O3 glass is the rate limiting step. Above 1200 C, the rate constants exhibit no dependence on P o2 , which is expected if oxygen transport through ZrO2 is the rate limiting step [79]. The addition of SiC has the effect of increasing the temperature for which glass transport properties provide the rate limiting step. The role of oxygen transport skeleton in ZrB2 + SiC materials is not clear at through the crystalline ZrO2 this time. Compositional variations, which form glasses that are stable to higher temperatures, and have low O2 diffusion rates and low vapor pressures, could further improve the oxidation resistance of these materials. Glasses have been identified which are composed of refractory oxides (La2O3 , ZrO2 , ThO2 , Ta2O5 , TiO2 , WO3 , and B2O3 ) [85]. It is recommended that new UHTC materials which exploit these glass-forming compositions be investigated. Compositional variations, which reduce oxygen transport through the skeleton phase, may improve the oxidation resistance as well. It is known that additions of higher valence metals into the ZrO2 lattice will reduce oxygen vacancy concentration and diffusion [79]. The pyrochlore phases associated with ZrO2 and rare earth oxides are also known to exhibit significantly lower oxygen diffusion than ZrO2 , as shown in Figs 12 and 13 [86, 87]. The in-situ formation of pyrochlore phases in UHTCs, and subsequent oxidation behavior, has not been investigated. It is recommended that new UHTC materials with rare earth additions be investigated. Vapor pressures at the metal-metal oxide interface are compared in Table I for potential skeleton phase-formers including Zr, Hf, Sc, Y, Be, and Ta. Materials forming HfO2 and ZrO2 are the best materials from this interfacial vapor pressure criterion. The carbides and nitrides of Zr (and Hf) are inferior to the diborides in oxidation behavior. Arc-heater testing of the carbides of Zr and Hf has been conducted at temperatures of 2400 to 2700 C with exposure times of 30 to 180 s in a gas environment replicating stoichiometric     Figure 12 Oxygen diffusivity for ZrO2 -rare earth oxide ceramics.  ULTRA-HIGH TEMPERATURE CERAMICS  Figure 13 Oxygen permeability constants for oxide ceramics with py rochlore structure.           hydrocarbon combustion [88]. A polished cross-section of HfC oxidized in arc-heater testing at 2700 for 30 s is shown in Fig. 14, revealing the porous oxide scale [89]. Vapor pressures for CO formed at the ZrC-ZrO2 and HfC-HfO2 interfaces have been calculated and exceed 1.01 × 105 Pa at 1730 C for both systems [90]. Thus, the carbides of Zr and Hf are compromised by the high vapor pressures formed at the oxide-carbide interface. The formation of porous scales was similarly observed for the nitrides. The formation of Zr and Hf oxycarbide intermediate phases in the ZrC-ZrO2 and HfC-HfO2 systems have been reported, as the primary oxygen barriers [91, 92] and as stable phases in themselves [93, 94]. Hafnium carbide was reported as exhibiting superior oxidation resistance to HfB2 at 1400 to 2100 C [91], but is based on furnace testing which biases the results unfavorably against the borides, as discussed previously. Stable oxycarbide phases have been reported for some rare earth metals, which suggests that rare earth additions may be employed to stabilize the oxycarbide phase in the scale. It is suggested that scandium (or ScC1−x ) be investigated since Sc has the same atomic radius as Zr or Hf and should exhibit high solubility into the base materials ZrC or HfC. Carbon interlayers, forming between the oxide and the base material, have also been reported for these carbides [95, 96]. Finally, the carbides are vulnerable to forming a powdery and non-adherent oxide scale below approximately 1700 C, which has been attributed to an inability to sinter at these temperatures [29, 97, 98]. Sintered and zone-melted samples exhibited low-temperature     5895  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 14 Microstructure (polished cross-section) of HfC arc-heater oxidized at 2700  C for 30 s.     catastrophic oxidation [30], while chemically vapor deposited (CVD) materials have exhibited protective scale formation [91]. The nature of the low temperature oxidation vulnerability is not understood at this time.  5. Oxidation studies—Glass scale modification  Oxidation behavior of non-oxide ceramics depends highly on the properties of the oxidation product and on the combination of physical and chemical processes taking place on the surface exposed to oxygencontaining atmosphere. In general, the chemical composition and structure of an oxidized surface define the oxidation stability of a ceramic material. Modification of the chemical composition of the oxide surface layer, leading to decreased inward diffusion of oxygen, is one of the efficient ways of controlling oxidation resistance of non-oxide ceramics. This modification can be accomplished by changing the bulk composition, or the  5896  r     surface of ceramics using CVD, ion implantation, “pack cementation”, and other methods. For example, the oxidation resistance of ZrB2 ceramics was significantly improved by modifying the bulk composition with SiC leading to the formation of a protective surface layer of borosilicate glass during exposure to an oxygen-containing atmosphere [18, 20, 33, 34, 39, 84]. The research conducted at NSWCCD showed that a further improvement of the oxidation performance of ZrB2 -SiC ceramics (up to 1600 C to-date) can be accomplished by the addition of CrB2 , TiB2 , TaB2 , NbB2 , and VB2 [99]. Oxidation of the modifying diborides resulted in the formation of corresponding oxides in the surface borosilicate glass. It should be emphasized that the oxidation resistance of all of the modifying diborides (alone) is much lower than that of the ZrB2 and HfB2 ceramics [14, 22]. It is known from the literature that borate and silicate glasses containing Group IV-VI transition metal oxides show strong tendency to phase separation (immiscibility) [100]. Systems exhibiting immiscibility are characterized by steeply rising liquidus temperatures and increased viscosity. An increase in the viscosity decreases the oxygen diffusion rate through the oxide surface scale based on the Stokes—Einstein relationship [101], which shows that diffusivity is inversely proportional to viscosity. Another potential benefit of increased viscosity as well as increased liquidus temperature is the suppression of boria evaporation from the glass. The oxide effectiveness in enhancing immiscibility increases with increasing metallic element cation field strength, z /r 2 , where z is the valence and is the ionic radius [100-103]. Since the cation field strengths of Ti, Nb, Ta, Mo, Cr, and V are higher than that of Zr, these elements can be effective in promoting phase separation of the borosilicate glass formed on the surface of the ZrB2 -SiC ceramics. The concept of phase separation as a controlling factor in the oxidation protection of non-oxide ceramics is unique and has not been discussed in the literature. The modifying diboride additives were introduced into the ZrB2 -SiC ceramics in the amounts of 2- 20 mol% as a substitution for ZrB2 . The molar ratio of ZrB2 to SiC in all materials was maintained at 2 (25 vol% SiC). The ceramics were prepared by hot pressing starting mixtures, consisting of ZrB2 , SiC, and modifying additives, at 2100 C and 20 MPa for 0.5 h. Oxidation experiments were conducted by heating the samples in a furnace in air at 1200-1600 C, typically for 2 h. The samples were placed into the furnace at the test temperatures and then air quenched after the hold. Quenching of the samples was conducted to retain the high temperature condition of the surface layer for analysis. Additionally, the oxidation of the samples was characterized by thermal gravimetric analysis (TGA) during 5 h isothermal heating at different temperatures in the air-simulating oxygen/argon mixture. The composition and structure of the surface and crosssection of the oxidized ceramics were evaluated using X-ray diffraction (XRD), SEM, and EDS. Fig. 15 shows the results of isothermal (5 h) TGA heating at 1300 C of ZrB2 -SiC ceramics modified with           \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS     SEM studies showed the evidence of hightemperature phase separation on the surfaces of all the modified samples after both TGA and furnace oxidation tests. Fig. 16a shows the surface of the sample containing 10 mol% TaB2 after the TGA oxidation test. The presence of large (more than 100 µm) droplets of borosilicate glass periodically distributed in a partially crystallized glassy matrix is an indication of high temperature glass phase separation. The matrix glass is enriched with Zr and Ta (from the EDS analysis data). X-ray diffraction of the whole surface of the oxidized sample showed the presence of ZrO2 along with a small amount Ta Zr2.75O8 with melting temperature above 1700 C. The CrB2 -containing ceramics (Fig. 16b) exhibited multiple phase separation after furnace oxidation at 1500 C, with shells around chromium-rich droplets which is the consequence of incomplete diffusion during cooling. A periodic pattern of crystals on the surface of ceramics modified with NbB2 (Fig. 16c) implies the existence of glasses of different compositions at the test temperature. X-ray diffraction analysis of the surface showed that the crystalline phase is Nb2Zr6O17 with a melting temperature about 1500 C. The glassy phase contains only small amounts of Nb and Zr (from EDS analysis). The presence of glass immiscibility is also observed on the surface of the ceramics modified with 5 mol% VB2 after TGA oxidation at 1300 C (Fig. 16d). Circular microcracking around droplets in this sample           Figure 15  C of ZrB2 /SiC ceramics modified with 10 mol% CrB2 , TiB2 , NbB2 , VB2 , and TaB2 .  Isothermal TGA oxidation at 1300     10 mol% of CrB2 , NbB2 , TaB2 , TiB2 and VB2 . The presence of the corresponding oxides in the oxidized surface layer improved the oxidation performance of the baseline material. The lowest weight gain during oxidation was observed for the ceramics containing TaB2 . The thickness of the oxidized layer of this sample was less than half of that for the baseline material [99]. The weight gain decreased in the sequence of the modifying additives: CrB2 , TiB2 , NbB2 , VB2 , and TaB2 , correlat+3 , Ti +4 , ing well with the cation field strength for Cr +5 , V +4 , and Ta +5 [99]. Nb  Figure 16 SEM micrographs of the surface of oxidized ZrB2 -SiC ceramics modified with: (a) TaB2, (b) CrB2 , (c) NbB2 , and (d) VB2 .  5897  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 17  C of TiB2 and TiB2 -SiC ceramics modified with 20 mol% ZrB2 , CrB2 , NbB2 , and TaB2 .  Isothermal TGA oxidation at 1300     can be attributed to compositional differences resulting in thermal expansion coefficient mismatch and leading to the development of strain and cracking during cooling. X-ray diffraction showed the presence of V7O13 in addition to ZrO2 on the surface of the oxidized VB2 containing sample. The discovered correlation between oxidation resistance of ZrB2 and the presence of phase separation in the surface protective glass was successfully applied to improve the oxidation performance of CrB2 , TiB2 , TaB2 , NbB2 , ZrB2 -Si3N4 , Ti3SiC2, and Si3N4 ceramics. The major results for the TiB2 , Ti3SiC2 , and Si3N4 ceramics are discussed below. The oxidation behavior of the TiB2 ceramics modified with SiC and 5-20 mol% CrB2 , NbB2 , TaB2 , and ZrB2 was evaluated [104]. The ceramics were prepared by hot pressing starting mixtures consisting of TiB2 ,        SiC, and diboride modifiers at 2100 C and 20 MPa for 0.5 h. Both furnace heating and TGA were used to characterize the oxidation resistance of the ceramics. The data show that the TiB2 -SiC ceramics containing TaB2 , NbB2 , and CrB2 have the best oxidation performance at all temperatures and additive contents. Fig. 17 presents the results of isothermal TGA heating at 1300 C. The addition of ZrB2 led to the decrease in oxidation resistance, especially, at 20% loading. The effect of additives on the oxidation resistance of TiB2 ceramics correlates with the cation field strength values for the elements being highest +4 oxide, with cation field strength for Ta and Nb. The presence of Zr +4 , led to the low protecting capalower than that of Ti bilities of the surface glass. If the rule of mixtures, not cation field strength of elements, was a controlling factor in the properties of surface glass and its protective capabilities, the improvement in the oxidation behavior of TiB2 -SiC ceramics could be expected with the introduction of ZrB2 , having the highest oxidation resistance of all the studied diborides. The optical micrograph of the surface of the CrB2 -containing ceramics (Fig. 18) clearly shows the immiscibility of the glass with the Cr2O3 -rich green areas and TiO2 -rich brown areas. The Ti3SiC2 ceramics recently attracted considerable attention because of their unique microstructure resulting in the exceptional combination of properties such as a high melting point, high fracture toughness and thermal shock resistance, plasticity at elevated temperatures, high modulus, low hardness, easy machinability, and self-lubrication [105-107]. These properties make the Ti3SiC2 ceramics of very high practical importance in numerous applications. However, these ceramics experience significant oxidation at temperatures above 1000 C preventing their high-temperature application     Figure 18 Optical micrograph of the surface of TiB2 -SiC ceramics containing 20 mol% CrB2 after oxidation at 1200     C for 2 h.  5898  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  in oxidizing environments. The modification of the ceramics with 10 mol% of TaB2 , ZrB2 , CrB2 , NbB2 , and VB2 was undertaken to create phase-separated borosilicate glass on the surface of the ceramics and, thus, increase their oxidation resistance. The Ti3SiC2 ceramics were synthesized and densified by reactive hot-pressing at 1500 C and 20 MPa for 1 h in He using the novel displacement reaction:     Ti5Si3 + 4TiC + 2C → Ti3SiC2  The additives were introduced in the starting Ti5Si3 TiC-C mixtures. The samples were oxidized during air Figure 20 Oxidation of Ti3 SiC2 ceramics modified with 10 mol% TaB2 , ZrB2 , CrB2 , TiB2 , and NbB2 during furnace heating for 2 h.              furnace heating at 1000 to 1500 C, typically for two hours. The hot-pressed ceramics were fully dense and contained Ti3SiC2 , small amounts of TiC, and additional phases associated with the additives. The SEM micrographs (Fig. 19a) show the presence of high aspect ratio plate-like grains, which define the mechanical behavior of the ceramics [105-107]. The additives (TaB2 in Fig. 19b) did not noticeably change the microstructure of the baseline material. The unmodified ceramics started to oxidize substantially at 1100 C (Fig. 20). Of all the diboride additives, NbB2 and, especially, TaB2 significantly improved the oxidation behavior of the ceramics. After 2 h oxidation at 1400 C, the weight gain was 430 g/m2 for the baseline ceramics compared to 60 g/m2 for the ceramics containing 10 mol% TaB2 . After heating at 1300 C the thickness of the oxidized layer was about 250 and 50 µm for the baseline and 10 mol% TaB2—modified ceramics, respectively. The microstructure of the oxidized surfaces of the baseline and TaB2 -modified ceramics (Fig. 21) is notably different. A bimodal distribution of the TiO2 crystals with significantly different morphology and size is observed in the TaB2 -containing sample. The bimodal distribution and clustering of these crystals are probably an indication that crystallization occurred from immiscible glasses of different composition during hightemperature exposure. The pronounced effect of TaB2 and NbB2 forming corresponding oxides in the glass is related to the highest cation field strength of Nb and Ta compared to all other elements tested. Silicon nitride is one of the most promising candidates for high-temperature structural applications, such as hot section components for advanced gas turbines and high-efficiency microturbines. High-temperature applications of Si3N4 ceramics depend, to a very high degree, on their behavior in corrosive environments and, primarily, on their resistance to oxidation. The effect of transition-metal diborides (CrB2 , TaB2 , and ZrB2 ) was studied on the oxidation behavior and microstructure of oxidized Si3N4 ceramics containing 5 wt% Y2O3 and 2 wt% Al2O3 as sintering aids [108]. The ceramics were hot-pressed at 1825 C and 20 MPa  Figure 19 SEM micrographs of the fracture surface of Ti3 SiC2 ceramics: (a) baseline Ti3 SiC2 and (b) Ti3 SiC2 modified with 10 mol% TaB2 .     5899  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 21 SEM micrographs of the surface of Ti3 SiC2 ceramics after oxidation at 1200 with 10 mol% TaB2 .     C for 2 h: (a) baseline Ti3 SiC2 and (b) Ti3 SiC2 modified  5900  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS           in He for 1 h. The oxidation behavior was characterized after furnace heating at 1200-1600 C in air. Fig. 22 shows that no significant oxidation (weight change) was observed for any of the materials below 1300 C. However, above 1350 C, only CrB2 significantly increased the oxidation resistance of the Si3N4 ceramics. The effect of the additives increases with increasing temperature. The SEM observations showed that the oxidized surface of the CrB2 -modified ceramics is almost fully crystallized, while the surface of the baseline ceramics is covered by a poorly crystallized, phase-separated glass (Fig. 23). The XRD analysis identified α -cristobalite as the only crystalline phase on the surface of the baseline material. In contrast, yttrium disilicate (monoclinic β -Y2Si2O7 ) together with a small amount of α -cristobalite were found on the surface of the Cr-containing material. No Cr-containing  Figure 22 Oxidation of Si3N4 ceramics modified with 10 mol% ZrB2 , TaB2 , and CrB2 .  Figure 23 SEM micrographs of the surface of Si3N4 ceramics after oxidation at 1550 vol% CrB2 .     C for 2 h: (a) baseline Si3N4 and (b) Si3N4 modified with 5  5901  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS     compounds were detected, however, on the surface of the materials by either EDS or XRD. It is hypothesized that chromium oxide present in the glass during oxidation additionally contributes to phase separation in the glass with the formation of a B-O-Si droplet phase and a Al-Y-Cr-Si-B-O matrix phase. Homogeneous nucleation and the formation of Cr2O3 crystallization centers occur in the matrix phase, followed by the catalytic crystallization of elongated Y2Si2O7 grains, while a small amount of fine cristobalite grains tration of Y2O3 · 2SiO2 crystals (melting temperature crystallizes from the droplet phase. The high concen1775 C) on the surface of the Cr-containing ceramics provides effective oxidation protection during exposure to oxidizing atmosphere. The highest oxidation resistance was shown by the ceramics containing less than 5 vol% CrB2 . The results of this research clearly showed that the oxidation resistance of non-oxide ceramics can be significantly enhanced by compositional design leading to the formation of a surface layer of immiscible multicomponent glass. The resulting increased liquidus temperatures and viscosities, as well as decreased oxygen diffusivities, in the immiscible glasses are considered responsible for the observed improvement in the oxidation resistance of the ceramics. The difference in the immiscibility and the corresponding oxidation behavior for different ceramics and modifying additives is a function of the oxidation state of elements in a particular glass, which highly affects their cation field strength and, consequently, the tendency of the glass toward phase separation. For example, Cr exhibited different oxidation states depending on the base ceramic material, and test temperature and atmosphere (TGA versus furnace oxidation). The presence of CrO2 was detected in the oxidation layer of several samples. The effective cation field strength of +4 is much higher than Cr +3 (1322 compared to 793 Cr −2 ) while that of Ta +5 is constant at 1220 nm −2 . This nm offers an explanation for the alternating effectiveness of Ta and Cr modifiers in enhancing oxidation resistance.  6. Conclusions     High vapor pressures at the metal-metal oxide interfaces of the slow-growing oxides (SiO2 , Al2O3 , Cr2O3 , BeO) are disruptive at 1800 C or below. The high interfacial vapor pressures of these systems primarily result from the system thermochemistry and are only secondarily dependent on the oxidation kinetics. Among these oxides, a glass-forming SiO2 scale exhibits significantly higher temperature capability, compared to materials which form a crystalline Al2O3 , Cr2O3 , or BeO scale, due to the greater structural tolerance of glass to high interfacial vapor pressures. Materials that form a multi-component oxide scale, composed of a refractory oxide skeleton and an amorphous (glass) oxide component, provide good oxidation performance at hypersonic use temperatures up to, and above, 2000 C. This multi-component oxide system is the only structure known at this time that mitigates, or recovers from, high interfacial vapor pressures. The     5902     oxidation resistance of ZrB2 -SiC and other non-oxide materials is improved, to at least 1600 C, by compositional modifications with transition metal additives that promote immiscibility in the glass component of the scale. The oxidation mechanisms of materials forming this scale structure (e.g., ZrB2 -SiC) are still not wellunderstood despite 40 years of research.  Acknowledgements  Sustained support by Dr. Stephen Fishman of the Office of Naval Research, and by the Naval Surface Warfare Center (Dahlgren and Carderock Divisions), is gratefully acknowledged. The authors would like also to thank Dr. E. 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Chem. 19 (1993) 1.  5903  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  104.  J .  A .  Z A Y K O S K I  et al., FY 2000 NSWC Res. Digest  (2000)  108.  I .  G .  T A L M Y et al.,  in Proceedings of the International Sym posium on “High Temperature Corrosion and Materials Chem istry IV,” edited by E. Opila, P. Hou, T. Maruyama, D. Shifler and  E. Wuchina (The Electrochemical Society, Inc., Pennington, NJ,  95.  105. M . W . B A R S O U M and T . E L R A G H Y , J. Amer. Ceram. Soc.  79 (1996) 1953.  J .  F R Y T and L .  106.  S T O B I E R S K I , Trans Tech Pubs., Switzerland,  2003) Vol. 2003-16, p. 361.  (1997) p. 1608. 107. N . (1999) 4385.  F . G A O , M I Y A M O T O and D .  Z H A N G , J. Mater. Sci. 18  Received 28 February and accepted 29 April 2004  5904  \\x0c']"
},{
  "_id": 200,
  "PDF": "Oxidation-behaviors-of-HfZrTaNbC-and-HfZrTaNbCSiC-at-13001500-C2021Journal-of-Materials-Science-and-Technology.pdf",
  "Text": "['Journal of Materials Science & Technology 60 (2021) 147-155  Contents lists available at ScienceDirect  Journal of Materials Science & Technology  j o u r n a l h o m e p a g e : w w w . j m s t . o r g  Research Article  Oxidation behaviors of (Hf0.25Zr0.25Ta0.25Nb0.25)C and (Hf0.25Zr0.25Ta0.25Nb0.25)C-SiC at 1300-1500  C  Haoxuan Wang a , Shouye Wang a , Yejie Cao a,∗ , Wen Liu b , Yiguang Wang c,∗  a Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China b School of Materials Science and Engineering, Zhengzhou University, Zhengzhou, Henan 450052, PR China c Institute of Advanced Structure Technology, Beijing Institute of Technology, Haidian District 100081, Beijing, PR China  a r t   i c l e   i n f o  a b s t r a c t  Article history:  Received 7 March 2020 Received in revised form 21 May 2020 Accepted 21 May 2020 Available online 7 July 2020  Keywords:  High-entropy carbide Oxidation resistance Oxidation mechanism  In this work, high-entropy ceramics (Hf0.25 Zr0.25 Ta0.25Nb0.25 )C (HZTNC) and HZTNC doped with 20 vol% SiC (HZTNC-SiC) were fabricated by spark plasma sintering. Their oxidation behavior was  investigated over the temperature range of 1300-1500   C for up to 60 min. Both HZTNC and HZTNC-SiC exhibited good oxidation resistance, and their weight change as a function of oxidation time obeyed a parabolic law. Through XRD, microstructure observation, and elemental mapping analysis of the oxide  layers,  it was found that the formation of Nb2 Zr6O17 , Hf6 Ta2O17 , and (Ta, Nb)2O5 mixed-oxide  layers effectively protected the matrix from further oxidation. Microcracks began to appear on the oxide layer of HZTNC at high temperatures after 60 min of oxidation. However, the addition of SiC in HZTNC suppressed these microcracks at high temperatures due to the active oxidation of SiC. Compared with the oxides formed on HZTNC,  the additional  formation of Hf(Zr)SiO4 on HZTNC-SiC could  further  improve  its oxidation resistance over HZTNC ceramics. © 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.  1.   Introduction  The unique properties of high-entropy alloys (HEAs) have stimulated  the  fast development of high-entropy ceramics  (HECs)  in recent years  [1-9]. HECs are single-phase compounds composed of cations or anions of  four or more  types of elements at either an equimolar or near-equimolar  ratio. The HECs,  like HEAs, are expected to exhibit  improved mechanical and  functional properties due to the “cocktail” effect [10-16]. In 2015, Sarkar et al. first developed  (MgNiZnCuCo)O anode materials with excellent storage capacity [17]. Since then, several HECs with different crystal structures have been synthesized, including materials with a rock salt  structure  [18-22], perovskite  structure  [23], fluorite  structure [24], spinel structure [25], AlB2 structure [1,26], half-Heusler structure [27], and CrSi2 structure [28]. These ceramics exhibited better mechanical properties  [29],  ionic conductivities  [30], corrosion resistance [31-35], nuclear waste immobilization [36], and low temperature conductivity [37] than the corresponding singlecomponent ceramics.  ∗ Corresponding authors. E-mail addresses: Caoyejie@nwpu.edu.cn (Y. Cao), wangyiguang@bit.edu.cn (Y. Wang).  Ultrahigh-temperature HECs (UHTHECs) have attracted particular attention due  to  their application  in areas such as  thermal protection systems of hypersonic vehicles [38], high-temperature devices [39], and radiation-resistant materials for advanced nuclear energy systems [40,41]. Ultrahigh-temperature ceramics are usually transition-metal carbides or borides, which are easily oxidized in oxygen-containing environments [42-45]. Thus, improving the oxidation resistance of these materials  is one of the key research topics  in  this field. Luo et al.  [42]  synthesized  the first boride UHTHEC, which had a significantly improved oxidation resistance compared with the corresponding single-component borides. Similarly, UHTHEC carbides also exhibited better oxidation resistance than the corresponding single-component carbides [35,42]. Since complex oxides are  formed during  the oxidation process due  to the multi-cation nature of the UHTHEC,  it can be argued that  its underlying oxidation mechanism  is controlled by oxygen  inward diffusion  [46] or  the cations outward diffusion  [47]  through  the oxide layer. However, it is generally believed that the formation of complex oxides during the oxidation of UHTHECs and the subsequent hindered diffusion effect of high-entropy cations account for the improvement in oxidation resistance [31,35]. In our previous study [35], a (Hf0.2 Zr0.2 Ti0.2 Ta0.2Nb0.2 )C (HZTTNC) UHTHEC was prepared with  improved oxidation resistance compared  to  the  individual  carbides, which was  still  inferior  https://doi.org/10.1016/j.jmst.2020.05.037 1005-0302/© 2020 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.  \\x0c', '148   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  compared  to  that of UHTCs with  the addition of SiC  (such as ZrB2 -SiC) [47-49]. When SiC was added to HZTTNC material [31], the resultant ceramics (HZTTNC-SiC) showed  improved oxidation resistance, albeit still not as good as that of ZrB2 -SiC. The outward diffusion of TiO was proposed to be the dominant process  in the oxidation of both HZTTNC [35] and HZTTNC-SiC [31]. Due to this process, the addition of SiC in HZTTNC did not lead to an improvement  in the oxidation resistance of HZTTNC compared to that of ZrB2 -SiC ceramics. With the aim of  further  improving the oxidation  resistance, TiC was  removed  from HZTTNC  to  form  (HZTNC) materials  in  this work,  thus avoiding the formation of TiO during oxidation. HZTNC doped with 20 vol% SiC (HZTNC-SiC) was also prepared. The oxidation behavior of the prepared materials was assessed  in the temperature range of 1300-1500  C to evaluate whether this strategy could  further enhance the oxidation resistance of UHTHECs.  It  is expected that this research would be helpful to  improve the high-temperature properties of UHTHECs by adjusting their composition.  Zr0.25 Ta0.25Nb0.25 )C   (Hf0.25  2. Experimental methods  Mechanical alloying combined with spark plasma sintering was used to synthesize HZTNC and HZTNC-SiC. Four carbide powders of HfC, ZrC, TaC, and NbC (purchased from HWRK CHEM, Beijing, China; 99 %, average size: 2  \\u242em) were weighed at a 1:1:1:1 M ratio of HfC to ZrC to TaC to NbC. Ball-milling was performed in a planetary ball mill (PMDW, Nanjing, China) at a speed of 500 rpm for 18 h. In order to prevent the vials from overheating, the milling was stopped for 5 min every 15 min. The balling medium was tungsten carbide and alcohol was used as dispersant. After ball-milling, the obtained powder mixture was divided into two parts: one part was used for the sintering of HZTNC bulk and the other part was mixed with SiC powder (99 %, 2  \\u242em, HWRK CHEM, Beijing, China) at a volume fraction of 80:20. The powder mixed with SiC was placed in a nylon tank and ball-milled at 130 rpm for 24 h in a roller ball mill (GQM-2−5, TENCEN POWDER, Changsha, China). The as-milled powders (with and without SiC) were placed in a Ф20 graphite die  Fig. 1. XRD patterns of the as-received ceramics.  ×  ×  and sintered by spark plasma sintering (Beijing Ailin Furnace Technology Co, Ltd.) at 2000  C for 15 min under a uniaxial pressure of 35 MPa in vacuum at a heating rate of 150  C/min. The as-sintered HZTNC and HZTNC-SiC ceramics were cut into 10 mm   5 mm   5 mm bulk samples for oxidation tests. One side of the sample was polished to 1  \\u242em finish with diamond lapping films. The polished samples were then placed on hollow zirconia balls in a zirconia crucible with the polished side facing up. Isothermal oxidation was carried out in an alumina tube furnace (GSL-1600X, Hefei Kejing Materials Technology Co., Ltd., Hefei, China) at 1300-1500  C  in air. When  the preset  temperature was reached,  it was  initially maintained for 10 min and the furnace tube was evacuated. Then, the furnace tube was quickly filled with air within 30 s, and finally adjusted with flowing air at a rate of 10 mL/min. Following the completion of oxidation, the air was replaced with high-purity argon and the samples were cooled to room temperature. The sam Fig. 2. SEM images showing the polished surface of HZTNC and HZTNC-SiC ceramics (a) HZTNC, (b) HZTNCSiC.  \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   149  Fig. 3. SEM surface images of HZTNC and HZTNC-SiC ceramics after isothermal oxidation at different temperatures: (a) HZTNC, 1300   C, (b) HZTNC-SiC, 1300   C, (c) HZTNC, 1400   C, (d) HZTNC-SiC, 1400   C, (e) HZTNC, 1500   C, (f) HZTNC-SiC, 1500   C.  Table 1 Calculated parabolic rate constants and activation energies for HZTNC, HZTNC-SiC, HZTTNC [35], HZTTNC-SiC [31] and ZrB2 -SiC [46] ceramics.  kp, 1300   C  (mg2 /cm4 h   kp, 1400   C  (mg2 /cm4 h   kp, 1500   C  (mg2 /cm4 h   Activation energy (kJ/mol)  HZTNC  HZTNC-SiC  HZTTNC  HZTTNC-SiC  ZrB2 -SiC   26.31  6.36  53.90  18.30  11.81   47.46  15.03  81.76  32.67  -   64.07  22.42  152.10  56.69  38.91   103 146 134 130 212  ples were then weighed by an analytical balance with an accuracy of 0.01 mg (Mettler Tolendo AG135, Greifensee, Switzerland). The above oxidation procedure was repeated five  times  for  the  total oxidation time of 60 min at 1300  C, 1400  C, and 1500  C, respectively. All oxidized samples were embedded in epoxy resin and cut in the middle. The cross-section was then polished to 1  \\u242em finish with diamond lapping films. The HZTNC and HZTNC-SiC phases after  isothermal oxidation were characterized by X-ray diffraction (XRD; Rigaku D/max-2400, Tokyo, Japan) with Cu Ka radiation. The microstructure of the samples was characterized by scanning electron microscopy (SEM; JEOL 6700 F, Tokyo, Japan) together with energy dispersive spectroscopy (EDS) for elemental mapping.  3. Results and discussion  ±  ±  By using the Archimedes method, the relative densities of the HZTNC and HZTNC-SiC bulks were measured to be 9.85   0.27 (>98 % theoretical density) and 8.59   0.27 (>99 % theoretical density), respectively. The XRD patterns of HZTNC and HZTNC-SiC showed only one phase with FCC structure for HZTNC, similar to that of TaC, albeit with a slight shift of the diffraction peaks toward the right (Fig. 1). Compared to HZTTNC  in our previous study [35], despite the removal of TiC, the remaining  four components retained the single-phase HEC structure. The sharp peaks observed for HZTNCSiC were assigned to HZTNC and the weak peaks were attributed to SiC; no other peaks were observed in HZTNC-SiC (Fig. 1). The morphologies of the polished HZTNC and HZTNC-SiC surfaces (Fig. 2) were also characterized. The HZTNC ceramic showed uniform elemental distribution (Fig. 2a), while the HZTNC-SiC ceramic showed two regions, a gray region corresponding to the HZTNC phase and a dark region corresponding to the SiC phase (Fig. 2b). A few micropores were observed on the surface of HZTNC, whilst HZTNC-SiC was almost fully dense,  indicating that the addition of SiC helped to increase the relative density of the ceramics. Fig. 3 shows the morphologies of HZTNC and HZTNC-SiC samples after isothermal oxidation at 1300−1500  C for 60 min. It can be seen that the formed oxide layers were dense without obvious for both HZTNC and HZTNC-SiC oxidized at 1300  C and defects   \\x0c', '150   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 4.  (a-c) Square of the specific weight change as a  temperature.  function of oxidation time at 1300, 1400, and 1500   C;   (d) the relationship between   lnkp and the reciprocal of  1400  C. At 1500  C, a number of microcracks were observed on the surface of oxidized HZTNC, while the surface of HZTNC-SiC was still dense without obvious defects. The isothermal oxidation data of HZTNC and HZTNC-SiC at 1300  C, 1400  C, and 1500  C were plotted in the form of the square of the specific weight change as a function of oxidation time (Fig. 4). These data were fitted based on the following equation [50]:  w2=   kp t   (1)  t   where w  is  the specific weight change during oxidation,  is  the oxidation time, and kp is the parabolic oxidation rate. As seen from Fig. 4, these curves were  linearly fitted, showing that the oxidation process followed the parabolic rate law (Eq. (1)). The parabolic oxidation behavior indicates that the diffusion process dominated the oxidation of HZTNC and HZTNC-SiC. The kp values were calculated through the slope of the w2 -t plots and compared with the kp values for HZTTNC and HZTTNC-SiC [31,35] (Table 1).  Indeed, the for HZTNC at 1300-1500  C were  kp values  lower than those  for HZTTNC, indicating that the removal of TiC led to an improvement in the oxidation resistance of UHTHECs. HZTNC-SiC also showed better oxidation resistance than HZTTNC-SiC, and even superior to that of ZrB2 -SiC ceramics [44-49]. The  increase  in kp value with  the  increase  in oxidation  temperature  indicates  the  faster  oxidation  of  ceramics  at  higher temperatures. The temperature dependence of kp can be described by Arrhenius equation, as follows [50]:  kp=   k0 exp(-E/RT)   (2)  where k0 is a pre-exponential constant, E is the oxidation activation energy, R  is the gas constant, and T  is the oxidation temperature. The natural logarithm of kp at 1300-1500  C (ln kp ) was plotted as a function of reciprocal temperature (1/T) in Fig. 4d. The oxidation  activation energies of HZTNC and HZTNC-SiC were calculated from the slopes of the curves in Fig. 4d and listed in Table 1. The oxidation activation energies were 103 kJ/mol for HZTNC, 146 kJ/mol for HZTNC-SiC, and 134 kJ/mol for HZTTNC [31,35]. The TiO outward diffusion during the oxidation of HZTTNC [35] cannot be the dominant step  in the oxidation mechanism of HZTNC and HZTNC-SiC due to the absence of TiC in both ceramics. Additionally, given their different oxidation activation energies, the controlling step in the oxidation of HZTNC would also be different from that in HZTNC-SiC. The outer oxide layers of HZTNC and HZTNC-SiC ceramics oxidized at 1300-1500  C  for 60 min were characterized by XRD (Fig. 5). The main phases of the outer oxide  layer of HZTNC were Nb2 Zr6O17 and Hf6 Ta2O17 , with a  small amount of  (Nb,Ta)2O5 (Fig. 5a). Without  the outward diffusion and evaporation of TiO, like  in HZTTNC,  the outer  layer of Nb2 Zr6O17 , Hf6 Ta2O17 ,  and (Nb,Ta)2O5 oxides mixture was dense, providing improved oxidation protection for HZTNC compared to HZTTNC. With the increase in oxidation temperature, the concentration of (Nb,Ta)2O5 (Fig. 5a) was  reduced,  indicating  that  the  formation  rate of  (Nb,Ta)2O5 was  inhibited at higher  temperatures. Thus, Ta and Nb  tended to  form Nb2 Zr6O17 and Hf6 Ta2O17 .  In addition  to  the oxidation products of HZTNC  in  the oxidized HZTNC-SiC  samples, a new phase of (Hf,Zr)SiO4 was also observed (Fig. 5b). The solid solution (Hf,Zr)SiO4 was  likely derived  from  the  reactions of  the oxidation products of HfC, ZrC, and SiC  in HZTNC-SiC. The presence of (Hf,Zr)SiO4 in the oxide layer enhanced the oxidation resistance of HZTNC-SiC. The cross-sectional morphologies and  the corresponding element mapping of the HZTNC and HZTNC-SiC ceramics oxidized at 1300  C  for 20 min and 60 min were analyzed (Figs. 6 and 7). A dense oxide layer was formed on the top of HZTNC or HZTNC-SiC ceramics after both 20 min and 60 min. Based on the elemental  \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   151  Fig. 5. XRD patterns of oxidized HZTNC and HZTNC-SiC ceramics at different temperatures: (a) HZTNC, (b) HZTNC-SiC.  Fig. 6. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1300   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  \\x0c', '152   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 7. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNCSiC ceramics at 1300   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min.  Fig. 8. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1400   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  mapping analysis, Hf, Zr, Ta, and Nb were found to be distributed uniformly in the oxide layer of HZTNC ceramics. Thus, according to the XRD results, it was hypothesized that HfO2 reacts with Ta2O5 to form Hf6 Ta2O17 , and ZrO2 reacts with Nb2O5 to form Nb2 Zr6O17 in the oxide  layer. Subsequently, the unreacted Ta2O5 and Nb2O5 form the solid solution of Nb(Ta)2O5 . As seen from the elemental mapping results of the oxidized HZTNC-SiC ceramics, the elements of Zr, Hf, and Si were enriched in the oxide layer while Ta and Nb  were also detected  in the oxide  layer. This result  is  in agreement with the XRD analysis which showed that the oxide phases were (Hf,Zr)SiO4 , Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 . At 1400  C, the oxide layers on HZTNC and HZTNC-SiC increased in thickness compared to those at 1300  C (Figs. 8 and 9), whereas the elemental distribution and oxide structure were similar to those at 1300  C. The cross-sectional morphologies of  the HZTNC and HZTNC C were  samples  oxidized  at  1500 observed  by  SEM  SiC   \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   153  Fig. 9. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNCSiC ceramics at 1400   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min.  Fig. 10. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC ceramics at 1500   C for different time: (a) HZTNC, 20 min, (b) HZTNC, 60 min.  (Figs. 10 and 11). The oxide  layer of HZTNC at 20 min was dense with uniformly distributed elements. Following 60 min of oxidation time, two oxide  layers were observed on the HZTNC surface: a Zrand Hf-enriched top layer with some lateral microcracks and pores, and a relatively dense second layer formed by a mixture of Hf, Zr, Ta, and Nb oxides. Due to the absence of TiO diffusion during oxidation, the formed oxide layer of HZTNC was dense after the first 20 min. With the increase in oxidation time at 1500  C, the thick ness of  the dense  layer  increased. Generally, stress  is developed with the growth in the oxide layer due to the difference in properties of the matrix and the oxides on  its top [51-53]. It  is believed that the stress in the dense oxide layer of HZTNC increases with the increase  in thickness, finally becoming  large enough to generate microcracks in the oxide layer. These microcracks provide diffusion channels for oxygen to further oxidize HZTNC to form a dense second oxide layer. For HZTNC-SiC ceramics, the oxide layers formed  \\x0c', '154   H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155  Fig. 11. Cross-sectional morphologies and EDX mapping for five constituting elements of the oxidized HZTNC-SiC ceramics at 1500   C for different time: (a) HZTNC-SiC, 20 min, (b) HZTNC-SiC, 60 min. (*1 Hf(Zr)TiO4 , Hf(Zr)SiO4 , Nb(Ta)2 O5 layer; 2 Hf-Zr-Ta-Nb-Si-O layer).  Fig. 12. Schematic oxide structure of HZTNC and HZTNC-SiC ceramics: (a) HZTNC, (b) HZTNC-SiC.  at 20 min and 60 min were both dense without obvious cracks. When the ceramic was oxidized for 60 min, two oxide layers were observed. The outermost dense layer was mainly composed of Zr, Hf, and Si, corresponding to (Hf,Zr)SiO4 solid solution according to XRD analysis. The second layer consisted primarily of Hf, Zr, Ta, and Nb oxides based on elemental analysis, and was depleted in the Si element. Similar phenomena were also observed  in the oxidation of ZrB2 -SiC ceramics [44,46-49,54], in which the Si-depleted layer was detected. Therefore, it is proposed that SiO outward diffusion occurred during the oxidation of HZTNC-SiC ceramics. The outward diffusion and evaporation of SiO may release the growth stress and avoid the formation of microcracks as the oxide layer thickens. Based on the above discussion, the plausible oxidation mechanisms for HZTNC and HZTNC-SiC are proposed, as shown in Fig. 12. According to the oxidation kinetics, the diffusion process should be dominant during the oxidation of HZTNC. It is unclear whether the controlling process is the inward oxygen diffusion or the outward  diffusion of cations such as Hf/Zr. However, the oxidation activation energy of HZTNC is only 103 kJ/mol, much lower than that required for cation outward diffusion [35]. Therefore, the  inward diffusion of oxygen through the  formed oxide  layer  is  likely the dominant process  for the oxidation of HZTNC  (Fig. 12a). During the oxidation process, a dense complex oxide layer composed of Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 is  formed on HZTNC, preventing  the matrix  from  further oxidation. As the dense  layer grows, growth stress accumulates and becomes  large enough to break the  layer and  form microcracks. A new, dense oxide  layer  is  then  formed below the  former  layer by the diffusion of oxygen through these cracks for further oxidation. The formation of the Si-depleted layer indicates that the oxidation mechanism of HZTNC-SiC is similar to that of ZrB2 -SiC, wherein the outward diffusion of SiO  is the controlling step during oxidation [44,54]. At low temperatures, the diffusion of SiO is so slow that the Si-depleted layer is not obvious and only a dense layer can be  \\x0c', 'H. Wang et al. / Journal of Materials Science & Technology 60 (2021) 147-155   155  observed. At high temperatures and long oxidation times, the diffusion of SiO is accelerated and the Si-depleted layer becomes more obvious. The diffused SiO reacts with the Zr/Hf oxides to form an additional dense  layer of (Hf, Zr)SiO4 , thus preventing the matrix from oxidation. Moreover, the outward diffusion of SiO also releases oxide growth stress, ensuring the integrity of the protective oxide layer. The oxidation resistance of HZTNC-SiC is thereby enhanced compared to that of ZrB2 -SiC ceramics.  4. Conclusion  In conclusion, the oxidation behaviors of HZTNC and HZTNC-SiC ceramics were assessed over the temperature range of 1300-1500  C in air for 1 h. The results indicated that both HZTNC and HZTNCSiC ceramics exhibited parabolic oxidation behavior over the tested temperature range. The inward diffusion of oxygen was proposed as  the possible oxidation mechanism  for HZTNC ceramics,  leading to the  formation of a dense mixed oxide  layer of Hf6 Ta2O17 , Nb2 Zr6O17 , and Nb(Ta)2O5 . The accumulated growth stress could break the oxide layer at high temperatures and long oxidation time, resulting in cracks and further oxidation of the matrix below. After the addition of SiC in HZTNC, the outward diffusion of SiO became the controlling process. This further improved the oxidation resistance of the ceramics and released the growth stress, thus avoiding the generation of microcracks. The multiple components of HECs can synergistically affect the properties of the materials. The present study provides an example of how the oxidation resistance of UHTHECs can be  improved by removing TiC  from HZTTNC to  form a new HEC, thus demonstrating that the properties of HECs can be adjusted by tuning their composition.  Acknowledgements  This work was financially supported by the National Natural Science Foundation of China (Grant No. 51972027). We would like to thank the Analytical & Testing Center of Northwestern Polytechnical University for assistance with SEM observations.  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},{
  "_id": 202,
  "PDF": "Oxidation-resistant ZrB2–SiC composites at 2200 C.pdf",
  "Text": "['Oxidation-resistant ZrB2-SiC composites at 2200 °C  Jiecai Han, Ping Hu, Xinghong Zhang *, Songhe Meng, Wenbo Han  Center for Composites Materials, Harbin Institute of Technology, Harbin 150001, PR China  Received 28 April 2007; received in revised form 16 August 2007; accepted 23 August 2007  Available online 1 September 2007  Abstract  Oxidation resistance tests were carried out on hot-pressed ZrB2-20 vol%SiC using an oxyacetylene torch. The temperature of the oxidized specimens exceeded 2200 °C. The mass and linear oxidation rates of the ZrB2-20 vol%SiC composites for 10 min were \\x000.23 mg/s and 0.66 lm/s, respectively. The surface layer appeared dense and adherent with the exception of a few burst bubbles and craterlets. No  macro-cracks or  spallation were detected after oxidation,  suggesting that  these composites possess a super oxidation resistance. The  microstructure of  the surface and cross-section of  the oxidized specimens were studied by scanning electron microscopy along with  energy dispersive spectrometry and X-ray diffraction. The oxidation mechanism was also discussed.  Ó 2007 Elsevier Ltd. All rights reserved.  Keywords: A. Ultra-high-temperature ceramic (UHTC); A. ZrB2; A. SiC; B. Oxidation resistance  1. Introduction  Ultra-high-temperature  ceramics  (UHTCs)  including  some  refractory metal  diborides  have  been  historically  studied and developed since 1960s [1]. These materials are  potential  candidates  for  thermal protection materials  in  both  re-entry  and  hypersonic  vehicles  because  of  their  excellent and unique combination of high melting points,  good thermal-shock resistance and good ablation/oxida tion resistance [2-7]. These properties make UHTCs attrac tive for  the design of  future hypersonic aerospace vehicle  with features  like  sharp leading  edges  and sharp nose cones. Such design features could produce more agile vehi cles that would open up a greater range of hypersonic flight  paths and re-entry trajectories. Re-entry and hypersonic  vehicles, regardless of their specific designs, require control  surfaces with sharp leading edges if they are to be maneu verable at hypersonic velocities. Low-radius leading edges  are subject to much greater aerothermal heating than blunt  edges, such as those on the space shuttle orbiter, and these edges will reach temperatures that may exceed 2000 °C dur ing re-entry [5,8]. The currently available thermal protec tion  materials  will  not  survive  under  such  extreme  temperatures  and new materials  are  required for use  in  advanced thermal protection systems.  ZrB2-SiC and HfB2-SiC are baseline UHTCs. Indeed, varying the starting composition  currently considered the  of  these UHTCs by changing the SiC content has given  added  flexibility  in  optimizing  specific microstructural  designs: adjusting the SiC content in ZrB2 and HfB2 matrices, for instance, has proven beneficial for improving the  oxidation and ablation resistance, without being detrimen tal to the high temperature stability [6,9]. For temperatures above 1200 °C, the addition of SiC provides more efficient  oxidation resistance by encouraging the formation of boro silicate glass on the  exposed surfaces, providing a much  greater oxidation protection than B2O3 alone [6,7,10,11]. Chemical interactions between the reaction scale formed  during oxidation and the phases present  in the bulk are  very complex and strongly influence the oxidation behav ior. The growth of  large bubbles underneath the external  reaction scale often causes  local  spalling which,  in turn,  noticeably abates the mechanical properties.  In the future,  the employment of UHTC is expected to  have a broad impact on several space applications, ranging  0266-3538/$ see front matter Ó 2007 Elsevier Ltd. All rights reserved.  doi:10.1016/j.compscitech.2007.08.017  * Corresponding author. Tel./fax: +86 45186402382.  E-mail address: zhangxh@hit.edu.cn (X. Zhang).  www.elsevier.com/locate/compscitech  Available online at www.sciencedirect.com  Composites Science and Technology 68 (2008) 799-806  COMPOSITES SCIENCE AND TECHNOLOGY  \\x0c', 'from thermal protection structures (TPS) for modern space  vehicles to propulsion for rocket motors or jet engines [4,8].  Concrete perspectives are  currently being addressed with  regards  to the re-design of a new generation of  re-usable  space vehicles with very sharp (not actively-cooled) leading  edges [12]. This innovative class of materials has the poten tial  to overturn the aerodynamic constraint that only TPS  with blunted profiles  can endure  the peak temperatures  generated when  a  vehicle  repeatedly  cuts  through  the  Earth’s atmospheric gases during hyper-sonic flights. The  highly-localized heating loads  foreseen in the  interaction  regions  dictate  the  need  for  outstanding  key  structural  issues  like oxidation resistance. The main techniques  for  the evaluation of the oxidation resistance include methods  involving  arc-jet,  furnace  and  oxyacetylene  torch  treat ment; however  the arc-jet experiment  is a very expensive  method, while  the  furnace  cannot achieve an ultra-high temperature oxidizing  environment. Simulated oxidation  tests on the ground are usually employed to study oxida tion properties in order to reduce the cost and give a pri mary  evaluation  of materials.  The  oxyacetylene  torch  method is the simplest technique to conduct with the lowest  cost. This can extend our knowledge of UHTCs by exam ining their oxidation behavior at a higher temperature than  can be obtained in a laboratory furnace.  The purpose of this paper is to evaluate the performance  and investigate the oxidation resistance and mechanism of  ZrB2-20%SiC at ultra high temperature using an oxyacetylene torch flame.  2. Experimental procedure  2.1. Preparation of ZrB2-SiC composites  Commercially available raw powders were used in this  study. The powders used for UHTC sample preparation  are  listed in Table  1. Prior  to hot pressing,  the ZrB2- 20 vol%SiC raw powders were milled in a polymer-coated  bucket charged with ethanol using WC balls  for 8 h and  dried in a rotating evaporator. Milled powders were uniax ially hot pressed in a boron nitride coated graphite die at 2000 °C  for  60 min  under  vacuum  (0.5 mbar),  and  30 MPa of applied pressure. Powder compacts were heated under vacuum (10\\x003 Pa) to 1700 °C with an average rate of \\x1820 °C/min. After the hot press was held at 1700 °C for it was backfilled with argon and heated at \\x1810 °C/ 0.5 h, 2000 °C. When to the die temperature reached  min  2000 °C, a uniaxial load of 30 MPa was applied. After the hot press was cooled at \\x1820 °C/min to room temperature. The load was removed when the die temperature dropped below 1750 °C. Four  60 min,  specimens  for  testing  were cut  from a billet with a diameter of 19 mm and an  overall height  of  14 mm. Prior  to oxidation,  specimens  were polished and cleaned in an ultrasonic bath in acetone,  and then weighed (accuracy of 0.1 mg). Models were placed  in graphite holders which enabled test durations in excess  of 10 min.  2.2. Oxidation equipment and testing  An oxyacetylene torch oxidation facility,  including a set  of temperature detection systems, a holder and an oxyacet ylene gun, was custom-built (see Fig. 1). The detecting sys tem includes  a  dynamic  responding multi-wavelength  pyrometer used for measuring the oxidation surface tem perature during oxidation. The specimen surface tempera ture can be controlled by regulating the distance between  the ablation gun and the specimen surface. The wavelength range of the pyrometer is 0.55-1.05 lm; O2, H2O, H+, OH\\x00 cannot the surface temperature detected by the pyrometer. Data  therefore, C2H2, interfere with the signals of  corresponding to the temperatures were recorded and pro cessed nominally every 1 s using a computer.  The oxidation behavior of ZrB2-SiC composites was tested with an oxyacetylene flame. The pressure and flux of acetylene were 0.1 MPa and 0.95 m3/h, and for oxygen and 1.90 m3/h,  0.5 MPa  respectively. The  specimen was  used for mass and linear ablation rate measurements.  The mass oxidation rate, Rm,  is defined as  Rm ¼ m1 \\x00 m2 t  ð1Þ  where m1 and m2 are the mass before and after oxidation, respectively, and t is the oxidation time.  The linear oxidation rate, Rd,  is defined as  Rd ¼ d t  ð2Þ  Table 1  Powders used for UHTC sample preparation  Material  Source  Mean  particle size  (lm)  Purity  (%)  ZrB2  Northwest Institute for Non-ferrous  Metal Research, China  5  99.5  SiC  Weifang Kaihua Micro-powder Co.,  Ltd., China  2  99.5  Fig. 1. Schematic of the ablation equipment, 1 - adjustable object carrier,  2 - graphite  cover, 3 - ablating sample, 4 - oxygen-ethyne flame, 5 -  ablating gun, 6 - support structure, 7 - multi-wavelength pyrometer, 8 -  computer.  800  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  \\x0c', 'where d and t represent the reaction layer thickness and the  3.2. Oxidation resistance of ZrB2-SiC  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  801  oxidation time, respectively.  The microstructures of  the oxidized surfaces and verti cal  sections of  the  specimens were  studied by  scanning  and after  Photographs of  the ZrB2-SiC composites before (right) oxyacetylene torch testing are shown in  (left)  electron microscopy (SEM; S-470, Hitachi,  Japan) along  Fig. 3. The surface layer appears dense and adherent with  with energy-dispersive spectroscopy (EDAX, USA) and (XRD; Rigaku, Dmax-rb, Cu Ka =  X-ray diffraction 1.5418 A˚ ).  3. Results and discussion  3.1. Density and microstructure  The bulk density of the hot-pressed sample was 5.41 g/ cm3, which corresponds to 98.2% of the theoretical density  the  exception of a few burst bubbles and craterlets. No  macro-cracks or spallation were detected. The temperature  of the oxidized surface of the ZrB2-SiC composites was in excess of 2200 °C, and the duration time was 10 min (see reaction layer was 375 lm.  Fig. 4). The  thickness of  the  The mass loss and linear oxidation rates of 2200 °C were  at  the ZrB2-SiC 0.23 mg/s and  composites for 10 min 0.66 lm/s, respectively.  Fig. 5 shows the surface morphology of at 2200 °C with different magnifications.  the oxide scale  It  is  significant  estimated with the rule of mixture. Fig. 2 shows scanning  to note that  the surface oxidized specimen consists of  two  electron micrographs of  the polished and fracture surface  distinctly different  structural  scales, namely,  smooth and  of  the ZrB2-SiC composites. The microstructure of composites is regular, in which the mean grain size was about 6 lm and little porosity was observed in the polished  the  rugged oxides. The fine oxide grains with size of almost 4 lm are not closely adherent  to each other with respect  to the first oxide (Fig. 5b), while the latter contain many  section. The SiC particulate, which represents the majority  micro-cracks and voids. From Fig. 5c and d,  it  can be  of the dark features in Fig. 2,  is homogenously distributed  found that the latter may have been molten under the pres in the diboride matrix and no agglomeration was detected.  The growth of ZrB2 grains was restricted due to the existence of SiC particles.  perature  ent test conditions. It should be noted that the testing tem500 °C lower (2680 °C)  temperature of pure ZrO2  [13]. The mechanism  melting  almost  than  the  is  Fig. 2. SEM micrographs of the polished and fracture surface of ZrB2-20%SiC.  Fig. 3. ZrB2-20%SiC specimen before (left) and after (right) oxyacetylene torch testing at 2200 °C for 10 min.  \\x0c', 'for the formation of this molten scale is most likely due to  the solution of SiO2 in ZrO2 point of the ZrO2 (Fig. 6). The ported to the surface and then deposited at outside scale.  resulting in a lower melting  liquid phase was  trans The burst bubbles  and craterlets upon the  surface was  found (see Figs. 3 and 5), which presumably derives from  the evolution of volatile products  like SiO(g), CO(g) and  B2O3(g) during oxidation. The EDS analysis of the surface of a specimen oxidized at 2200 °C for 10 min shows the  presence of zirconium and oxygen as the primary elements,  while silicon and boron are not detected by the apparatus.  XRD examination of  the oxide surface (not  shown)  indi cates  that  the  surface  layer was primarily  composed of  monoclinic ZrO2.  The  highlight  of  the  cross-section  of  the  specimen  exposed to an oxidizing environment  is given in Fig. 7.  The analysis  revealed the presence of five distinct  layers,  in contrast  to the three and four distinct  layers  reported  for Hf-based and Zr-based materials [3,14], which may be  caused by the different oxidation temperatures  since  the  temperature in the present case was much higher than that  have been reported for the most oxidation studies. The outlayer was \\x1870 lm thick and composed of porous ZrO2, as has been observed in Fig. 7. Beneath this layer was found a dense recrystallized layer \\x1880 lm thick (layer 2) that has not been reported for Hf-based and Zr-based  ermost  materials in the previous studies [3,14]. Further analysis of  layer 2 (Fig. 8) showed that it consisted of large oxide particles with the size of approximately 20 lm. The formation  of this layer is thought to be the consequence of the high oxi dation temperature since high temperature is beneficial for  the  sintering of  zirconium oxide and the  growth of  the  grains. The third layer  is composed of many voids which  are probably formed due  to the active oxidation of SiC  and formation of B2O3 with a high partial pressure. The voids appear to be on the order of 25 lm which have not  been observed in lower temperature testing. This is a transi tional region between the recrystallized oxide and the under lying  SiC-depletion  region  (layer  4). Apparently,  the  bonding strength at  the  interface between this  layer and  the SiC-depletion layer is weakest among all  layers due to  the existence of numerous  large voids and the coefficient  of the thermal expansion mismatch between the ZrO2 (10.8 · 10\\x006 K\\x001) [15] and unaltered ZrB2 (6.7 · 10\\x006 K\\x001) [16]. For weak bonding and high vapor pressure, cracks  are easily formed and spallation is liable to occur. Thus, this  layer plays a very important role in the performance of the  0  100  200  300  400  500  600  1000  1200  1400  1600  1800  2000  2200  2400  T  e  m  p  e  r  a  t  u  r  e  ˚ (  C  )  Time (s)  Fig.  4. Measured  results  of  temperature  curves  of  ZrB2-20%SiC  composites.  Fig. 5. Surface morphologies of the oxide scale after ZrB2-20%SiC specimen exposure to oxyacetylene torch testing at 2200 °C for 10 min.  802  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806    \\x0c', 'J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  803  Fig. 6. Phase diagram of the binary ZrO2-SiO2 system [13].  material at high temperature and the control of its behavior  by optimizing the ratio of ZrB2 to SiC becomes essential to the improvement of the oxidation property at high tempertemperatures above 2200 °C. Under ature, especially for  neath this oxide  scale  is  the SiC-depletion region, which  Fig. 8. SEM micrograph of the second layer in Fig. 7.  was similar to the high temperature oxidation behavior of  the ZrB2-SiC and HfB2-SiC composites reported in the previous studies [3,14]. According to the EDS analysis, this  layer is very poor in silicon, which indicates that once the sil icon carbide is oxidized, the Si diffuses directly towards the  outside scale. The SEM and EDS results indicated that  it  dominantly consisted of unaltered ZrB2, whose shape and structure were not changed significantly during the oxida tion processing and the 3D interconnected nature remained.  The fifth layer is the unaltered ZrB2-SiC. Oxidation of ZrB2-SiC at 2200 °C in the oxyacetylene torch testing suggests that ZrO2 becomes a protective layer  Fig. 7. SEM result for a cross-section of ZrB2-SiC after oxidation at 2200 °C for 10 min.  \\x0c', 'at these temperatures by recrystallizing into a dense coher ent  scale. The variation of  the  thickness of  the  reaction  layer with time can be an indicator of the oxidation mech anism. The parabolic trend (Fig. 9) suggests reaction diffu sion-controlled kinetics.  In this  case,  the dense  coherent  ZrO2 surface layer acted as a barrier for diffusion of oxygen to the underlying ZrB2-SiC.  3.3. Oxidation mechanism of ZrB2-SiC  There are three kinds of oxidants,  i.e., H2O, CO2 and case. H2O and CO2 were not taken into the sake of simplicity. Apparently,  O2, account  in present  for  the  exis tence of H2O will the SiO2 and B2O3 (i.e., HBO2, H3BO3, Si(OH)4, SiO(OH)2 and SiO(OH)) leading to increased oxidation rate. Therefore,  increase  the volatilization of  the main expected reactions during oxidation process are  as follows  ZrB2 ðsÞ þ 5 O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðlÞ B2O3 ðlÞ ! B2O3 ðgÞ 2  ð3Þ  ð4Þ  SiCðsÞ þ 3 O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ 2 SiO2 ðlÞ ! SiO2 ðgÞ SiO2 ðlÞ ! SiOðgÞ þ 1=2O2 ðgÞ  ð5Þ  ð6Þ ð7Þ ð8Þ  Oxidation resistance depends on a protective oxide scale  formed on the surface of  the matrixes. Whether the oxide  products can form a dense adherent scale is very important  for ultra-high-temperature applications. In some cases, this  depends on the gas pressure in the internal oxide scale. The  vapor pressures of  the predominant gases  for ZrB2-SiC composites were calculated to compile a volatility diagram at 2200 °C (Fig. 10).  In air, B2O3(g) vapor species and its pressure is 2.7 atm (reaction 4). How is  the predominant  ever, at low oxygen partial pressure, SiO(g) has the highest  vapor  pressure  up to  12.3 atm (reaction  6). They  both  exceed 1 atm.  If  the formed gaseous phase products with  high partial pressures  cannot  escape  rapidly through the  outside scale and the gas pressure exceeds ambient prescreated. At 2200 °C,  sure, voids would be  the oxidation  resistance of SiC is poor due  to active oxidation, which  results in the formation of high vapor pressure oxide prod ucts like SiO and CO with almost no silica, which provides  effective oxidation protection at elevated temperature. The  active oxidation of SiC is a key issue in the use of SiC con taining composites at extremely high temperatures.  The  thermodynamic  stability  diagram for ZrB2-SiC temperature range of 1500-  with different 2500 °C was  P O2  in  the  calculated (Fig.  11), designating  the  stable  phases in each region to interpret and analyze the oxidation  mechanism of this kind of ultra-high-temperature ceramics. At any temperature, ZrB2 and SiC are stable at P O2 values below both lines. In the range of investigated temperature, the equilibrium P O2 for ZrB2 is higher than SiC, that ZrB2 is more stable than SiC at ultra-high-temperature. The selective removal of SiC from this region suggests  implying  -40  -30  -20  -10  0  10  -30  -20  -10  0  10  ZrB2+SiO2  BO2  B2O3(g)  SiO(g)  ZrO2+SiO2  ZrB2+SiC  g o L  P  a  r  t  i  a  l  p  r  s s e  u  r  e  Log pO 2  Fig. 10. Volatility diagram for ZrB2-SiC at 2200 °C.  10  20  30  40  200  300  400  500  600  700  800  900  Y =48.75+36.375 X-0.4125 X2  h T  i  k c  n  s e s s e  o  f  t  h  e  r  c a e  t  i  n o  l  e y a  r  (  µ  m  )  Holding time (min)   Fig. 9. Variation of the reaction layer thickness with time for exposure of ZrB2-SiC at 2200 °C.  -18 1800  1900  2000  2100  2200  2300  2400  2500  -16  -14  -12  -10  -8  -6  -4  ZrB2 , SiO and CO  ZrO2 , B2O3 , SiO and CO  ZrB2 and SiC  g o L  P  O  2  (  a  t  m  )  Temperature (°C)   SiC+3/2O2  ZrB2+5/2O2  SiO2+CO ZrO2+B2O3  → →  Fig. 11. Thermodynamic P O2 at different  stability diagram for ZrB2-SiC with different  temperatures.  804  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806                \\x0c', 'that  the active oxidation of SiC occurred without  signifi cant oxidation of ZrB2. provided by the thermodynamic calculations (Fig. 11). At 2200 °C, the equilibrium for \\x1810\\x0010 atm while the equilibrium P O2 for ZrB2 oxidation is \\x1810\\x009 atm. During oxidation, the P O2 in the depletion layer, including the interface between the depletion layer  Justification for  this assertion is  P O2  SiC  oxidation  is  and the unaltered ZrB2-SiC would be set by the oxidation of SiC thereby protecting ZrB2 from attack. The main reaction products from the oxidation of ZrB2- SiC composites are ZrO2, SiO2 and B2O3. The performance of the ZrB2-SiC ultra-high-temperature ceramic and the determination of whether the oxide products can form a  stable oxide scale are important for ultra-high-temperature  applications. This  is  largely dependent on volatilization  and decomposition of the oxide products. The vapor pres sure vs.  temperature above some oxides at ambient pres sure was  calculated  as  shown  in Fig. 12. The vapor 2200 °C are 2.8 · 105 Pa  pressures of B2O3 and SiO2 and 18 Pa, respectively, whereas  at  the  vapor  pressure  of  ZrO2 In fact, both SiO2 and B2O3 have high vapor pressures, so at high temperature they both quickly  is 0.025 Pa.  vaporize.  In addition, SiO2 could decompose by another mechanism, such as reaction 8. Furthermore, the formation  of some gaseous oxidation by-products (i.e., CO(g)), espe cially at high partial pressures, would accelerate the volatil ization of  the initial  formed glass and create short-circuit  paths for the incoming oxygen resulting in increased oxida tion rate. The volatilization of ZrO2 was negligible comparing with the other two oxides in this case since its vapor  pressure  is  very  low,  almost  three orders of magnitude  lower  than SiO2. The mass active oxidation of SiC and the volatilization of the partial  loss  closely correlates  to the  oxidation products such as B2O3 and SiO2. The protective role of the oxide products  is different,  which was mostly determined by the temperature. At low temperature (below 1100 °C), ZrB2 oxidizes much more rapidly than SiC and the formed liquid B2O3 acts as a barrier to  oxygen diffusion resulting in passive oxidation of ZrB2-SiC composites. Above 1300 °C, B2O3 becomes non-protective because of rapid evaporation and the higher temperature  protection was ascribed to SiO2. SiC significantly improves the oxidation resistance of the composites due to the forma tion of the more viscous silica-containing glass, which has a  higher melting temperature, a lower vapor pressure, and is a  greater a barrier to oxygen diffusion. Moreover, SiC signifi cantly increases the temperature range of the glass as a pri mary oxygen barrier, and extends  the maximum working  temperature in oxidizing environments. In fact, SiC-contain ing ZrB2 based UHTCs are effective up to high temperatures of 1700-1800 °C. Nonetheless, at higher temperatures, SiC is  no longer responsible for the improvement in oxidation resis tance of the composites because of a lack of silica formation  during the oxidation process. On one side, SiC undergoes a  transition from passive to active oxidation where the protec tive SiO2 layer is removed as SiO and CO. On the other side, the initially formed silica-enriched glass will be lost by con siderable evaporation and decomposition (reactions (7) and (8)). For temperatures higher than 1800 °C, ZrO2 recrystallizes into a dense coherent subscale (Figs. 7 and 8). This  subscale composed of zirconia appears to be able to protect  the underlying ceramic from catastrophic oxidation.  4. Conclusions  The ZrB2-20 vol%SiC composites exhibited excellent oxidation resistance at 2200 °C. The thickness of the reaction layer is only 375 lm after 10 min oxidation. The mass  and linear oxidation rates of the ZrB2-SiC composites are \\x000.23 mg/s and 0.66 lm/s, respectively. No macro-cracks or spallation were detected after oxidation. SiC exhibited  preferential oxidation compared with ZrB2 study. The formation of high pressure gaseous  in the present  phases  (i.e., B2O3 and SiO) amount of pores. SiC was no longer  led to the generation of a significant  responsible  for  the  improvement  in oxidation resistance of  the composites  in  the present case. ZrO2 recrystallized into a dense coherent subscale, which protected the underlying ceramic from cat astrophic oxidation.  Acknowledgements  This work was supported by the National Natural Sci ence Foundation of China (90505015 and 50602010),  the  Research Fund for the Doctoral Program of Higher Edu cation (20060213031) and the Program for New Century  Excellent Talents in University.  References  [1] Kuriakose AK, Margrave JL. The oxidation kinetics of  zirconium  diboride and zirconium carbide at high temperatures. J Electrochem  Soc 1964;111(7):827-31.  [2] Monteverde F. The thermal  stability in air of hot-pressed diboride  matrix  composites  for uses  at ultra-high-temperatures. Corros Sci  2005;47(8):2020-33.  1000  1200  1400  1600  1800  2000  2200  2400  2600  10 -10  10 7 10 6 10 5 10 4 10 3 10 2 10 1 10 0 10 -1 10 -2 10 -3 10 -4 10 -5 10 -6 10 -7 10 -8 10 -9  ZrO2(g)  SiO2(g)  B2O3(g)  V  a  o p  r  p  r  s s e  u  r  e  (  P  a  )  Temperature (˚C)  Fig. 12. Vapor pressure vs. temperature above some oxides, calculated at  ambient pressure.  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  805      \\x0c', '806  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  [3] Chamberlain AL, Fahrenholtz WG, Hilmas G, Ellerby D. Oxidation  [10] Opila EJ, Halbig MC. Oxidation of ZrB2-SiC. Ceram Eng Sci Proc  of ZrB2-SiC ceramics under atmospheric and re-entry conditions.  2001;22(3):221-8.  Refract Appl Tran 2005;1(2):1-8.  [11] Rezaie Alireza, Fahrenholtz WG, Hilmas GE. Evolution of structure  [4] Chamberlain AL,  Fahrenholtz WG, Hilmas GE,  Ellerby DT.  Characterization  of  zirconium diboride  for  thermal  protection  during the oxidation of zirconium diboride-silicon carbide in air up to 1500 °C. J Eur Ceram Soc 2007;27:2495-501.  systems. Key Eng Mater 2004;264-268(I):493-6.  [5] Opeka MM, Talmy IG, Zaykoski JA. Oxidation-based materials selection for 2000 °C + hypersonic aerosurfaces:  theoretical  consid [12] Paul Kolodziej. Aerothermal performance  constraints  for hyperve locity  small  radius  unswept  leading  edges  and  nosetips. NASA  Technical Memorandum 112204; 1997.  erations and historical experience. J Mater Sci 2004;39(19):5887-904.  [13] Butterman WC, Foster WR. Zircon stability  and the ZrO2-SiO2  [6] Chamberlain AL, Fahrenholtz WG, Hilmas GE, Ellerby DT. High phase diagram. The Am Mineral 1967;52:880-5.  strength  zirconium diboride-based  ceramics.  J Am Ceram Soc  [14] Fahrenholtz WG. Thermodynamic analysis of ZrB2-SiC oxidation:  2004;87(6):1170-2.  formation  of  a  SiC-depletion  region.  J  Am  Ceram  Soc  [7] Monteverde F, Bellosi A. Oxidation of ZrB2-based ceramics  in dry  2007;90(1):143-8.  air. J Electrochem Soc 2003;150(11):B552-9.  [15] Li XT, Shi JL, Zhang GB, Zhang H, Guo QG, Liu L. Effect of ZrB2  [8] Van Wie DM, Drewry  Jr DG, King DE, Hudson CM.  The  on  the  ablation  properties  of  carbon  composites. Mater  Lett  hypersonic  environment:  required operating conditions and design  2006;60:892-6.  challenges. J Mater Sci 2004;39(19):5915-24.  [16] Loehman R, Corral E, Dumm HP, Kotula P, Tandon R. Ultra-high [9] Monteverde F, Bellosi A. Microstructure and properties of an HfB2-  temperature  ceramics  for  hypersonic  vehicle  applications.  Sandia  SiC composite  for ultrahigh-temperature  applications. Adv Eng  Report. 2006;1-46.  Mater 2004;6(5):331-6.  \\x0c']"
},{
  "_id": 203,
  "PDF": "Oxidation-resistant ZrB2–SiC composites at 2200°C.pdf",
  "Text": "['Oxidation-resistant ZrB2-SiC composites at 2200 °C  Jiecai Han, Ping Hu, Xinghong Zhang *, Songhe Meng, Wenbo Han  Center for Composites Materials, Harbin Institute of Technology, Harbin 150001, PR China  Received 28 April 2007; received in revised form 16 August 2007; accepted 23 August 2007  Available online 1 September 2007  Abstract  Oxidation resistance tests were carried out on hot-pressed ZrB2-20 vol%SiC using an oxyacetylene torch. The temperature of the oxidized specimens exceeded 2200 °C. The mass and linear oxidation rates of the ZrB2-20 vol%SiC composites for 10 min were \\x000.23 mg/s and 0.66 lm/s, respectively. The surface layer appeared dense and adherent with the exception of a few burst bubbles and craterlets. No  macro-cracks or  spallation were detected after oxidation,  suggesting that  these composites possess a super oxidation resistance. The  microstructure of  the surface and cross-section of  the oxidized specimens were studied by scanning electron microscopy along with  energy dispersive spectrometry and X-ray diffraction. The oxidation mechanism was also discussed.  Ó 2007 Elsevier Ltd. All rights reserved.  Keywords: A. Ultra-high-temperature ceramic (UHTC); A. ZrB2; A. SiC; B. Oxidation resistance  1. Introduction  Ultra-high-temperature  ceramics  (UHTCs)  including  some  refractory metal  diborides  have  been  historically  studied and developed since 1960s [1]. These materials are  potential  candidates  for  thermal protection materials  in  both  re-entry  and  hypersonic  vehicles  because  of  their  excellent and unique combination of high melting points,  good thermal-shock resistance and good ablation/oxida tion resistance [2-7]. These properties make UHTCs attrac tive for  the design of  future hypersonic aerospace vehicle  with features  like  sharp leading  edges  and sharp nose cones. Such design features could produce more agile vehi cles that would open up a greater range of hypersonic flight  paths and re-entry trajectories. Re-entry and hypersonic  vehicles, regardless of their specific designs, require control  surfaces with sharp leading edges if they are to be maneu verable at hypersonic velocities. Low-radius leading edges  are subject to much greater aerothermal heating than blunt  edges, such as those on the space shuttle orbiter, and these edges will reach temperatures that may exceed 2000 °C dur ing re-entry [5,8]. The currently available thermal protec tion  materials  will  not  survive  under  such  extreme  temperatures  and new materials  are  required for use  in  advanced thermal protection systems.  ZrB2-SiC and HfB2-SiC are baseline UHTCs. Indeed, varying the starting composition  currently considered the  of  these UHTCs by changing the SiC content has given  added  flexibility  in  optimizing  specific microstructural  designs: adjusting the SiC content in ZrB2 and HfB2 matrices, for instance, has proven beneficial for improving the  oxidation and ablation resistance, without being detrimen tal to the high temperature stability [6,9]. For temperatures above 1200 °C, the addition of SiC provides more efficient  oxidation resistance by encouraging the formation of boro silicate glass on the  exposed surfaces, providing a much  greater oxidation protection than B2O3 alone [6,7,10,11]. Chemical interactions between the reaction scale formed  during oxidation and the phases present  in the bulk are  very complex and strongly influence the oxidation behav ior. The growth of  large bubbles underneath the external  reaction scale often causes  local  spalling which,  in turn,  noticeably abates the mechanical properties.  In the future,  the employment of UHTC is expected to  have a broad impact on several space applications, ranging  0266-3538/$ see front matter Ó 2007 Elsevier Ltd. All rights reserved.  doi:10.1016/j.compscitech.2007.08.017  * Corresponding author. Tel./fax: +86 45186402382.  E-mail address: zhangxh@hit.edu.cn (X. Zhang).  www.elsevier.com/locate/compscitech  Available online at www.sciencedirect.com  Composites Science and Technology 68 (2008) 799-806  COMPOSITES SCIENCE AND TECHNOLOGY  \\x0c', 'from thermal protection structures (TPS) for modern space  vehicles to propulsion for rocket motors or jet engines [4,8].  Concrete perspectives are  currently being addressed with  regards  to the re-design of a new generation of  re-usable  space vehicles with very sharp (not actively-cooled) leading  edges [12]. This innovative class of materials has the poten tial  to overturn the aerodynamic constraint that only TPS  with blunted profiles  can endure  the peak temperatures  generated when  a  vehicle  repeatedly  cuts  through  the  Earth’s atmospheric gases during hyper-sonic flights. The  highly-localized heating loads  foreseen in the  interaction  regions  dictate  the  need  for  outstanding  key  structural  issues  like oxidation resistance. The main techniques  for  the evaluation of the oxidation resistance include methods  involving  arc-jet,  furnace  and  oxyacetylene  torch  treat ment; however  the arc-jet experiment  is a very expensive  method, while  the  furnace  cannot achieve an ultra-high temperature oxidizing  environment. Simulated oxidation  tests on the ground are usually employed to study oxida tion properties in order to reduce the cost and give a pri mary  evaluation  of materials.  The  oxyacetylene  torch  method is the simplest technique to conduct with the lowest  cost. This can extend our knowledge of UHTCs by exam ining their oxidation behavior at a higher temperature than  can be obtained in a laboratory furnace.  The purpose of this paper is to evaluate the performance  and investigate the oxidation resistance and mechanism of  ZrB2-20%SiC at ultra high temperature using an oxyacetylene torch flame.  2. Experimental procedure  2.1. Preparation of ZrB2-SiC composites  Commercially available raw powders were used in this  study. The powders used for UHTC sample preparation  are  listed in Table  1. Prior  to hot pressing,  the ZrB2- 20 vol%SiC raw powders were milled in a polymer-coated  bucket charged with ethanol using WC balls  for 8 h and  dried in a rotating evaporator. Milled powders were uniax ially hot pressed in a boron nitride coated graphite die at 2000 °C  for  60 min  under  vacuum  (0.5 mbar),  and  30 MPa of applied pressure. Powder compacts were heated under vacuum (10\\x003 Pa) to 1700 °C with an average rate of \\x1820 °C/min. After the hot press was held at 1700 °C for it was backfilled with argon and heated at \\x1810 °C/ 0.5 h, 2000 °C. When to the die temperature reached  min  2000 °C, a uniaxial load of 30 MPa was applied. After the hot press was cooled at \\x1820 °C/min to room temperature. The load was removed when the die temperature dropped below 1750 °C. Four  60 min,  specimens  for  testing  were cut  from a billet with a diameter of 19 mm and an  overall height  of  14 mm. Prior  to oxidation,  specimens  were polished and cleaned in an ultrasonic bath in acetone,  and then weighed (accuracy of 0.1 mg). Models were placed  in graphite holders which enabled test durations in excess  of 10 min.  2.2. Oxidation equipment and testing  An oxyacetylene torch oxidation facility,  including a set  of temperature detection systems, a holder and an oxyacet ylene gun, was custom-built (see Fig. 1). The detecting sys tem includes  a  dynamic  responding multi-wavelength  pyrometer used for measuring the oxidation surface tem perature during oxidation. The specimen surface tempera ture can be controlled by regulating the distance between  the ablation gun and the specimen surface. The wavelength range of the pyrometer is 0.55-1.05 lm; O2, H2O, H+, OH\\x00 cannot the surface temperature detected by the pyrometer. Data  therefore, C2H2, interfere with the signals of  corresponding to the temperatures were recorded and pro cessed nominally every 1 s using a computer.  The oxidation behavior of ZrB2-SiC composites was tested with an oxyacetylene flame. The pressure and flux of acetylene were 0.1 MPa and 0.95 m3/h, and for oxygen and 1.90 m3/h,  0.5 MPa  respectively. The  specimen was  used for mass and linear ablation rate measurements.  The mass oxidation rate, Rm,  is defined as  Rm ¼ m1 \\x00 m2 t  ð1Þ  where m1 and m2 are the mass before and after oxidation, respectively, and t is the oxidation time.  The linear oxidation rate, Rd,  is defined as  Rd ¼ d t  ð2Þ  Table 1  Powders used for UHTC sample preparation  Material  Source  Mean  particle size  (lm)  Purity  (%)  ZrB2  Northwest Institute for Non-ferrous  Metal Research, China  5  99.5  SiC  Weifang Kaihua Micro-powder Co.,  Ltd., China  2  99.5  Fig. 1. Schematic of the ablation equipment, 1 - adjustable object carrier,  2 - graphite  cover, 3 - ablating sample, 4 - oxygen-ethyne flame, 5 -  ablating gun, 6 - support structure, 7 - multi-wavelength pyrometer, 8 -  computer.  800  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  \\x0c', 'where d and t represent the reaction layer thickness and the  3.2. Oxidation resistance of ZrB2-SiC  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  801  oxidation time, respectively.  The microstructures of  the oxidized surfaces and verti cal  sections of  the  specimens were  studied by  scanning  and after  Photographs of  the ZrB2-SiC composites before (right) oxyacetylene torch testing are shown in  (left)  electron microscopy (SEM; S-470, Hitachi,  Japan) along  Fig. 3. The surface layer appears dense and adherent with  with energy-dispersive spectroscopy (EDAX, USA) and (XRD; Rigaku, Dmax-rb, Cu Ka =  X-ray diffraction 1.5418 A˚ ).  3. Results and discussion  3.1. Density and microstructure  The bulk density of the hot-pressed sample was 5.41 g/ cm3, which corresponds to 98.2% of the theoretical density  the  exception of a few burst bubbles and craterlets. No  macro-cracks or spallation were detected. The temperature  of the oxidized surface of the ZrB2-SiC composites was in excess of 2200 °C, and the duration time was 10 min (see reaction layer was 375 lm.  Fig. 4). The  thickness of  the  The mass loss and linear oxidation rates of 2200 °C were  at  the ZrB2-SiC 0.23 mg/s and  composites for 10 min 0.66 lm/s, respectively.  Fig. 5 shows the surface morphology of at 2200 °C with different magnifications.  the oxide scale  It  is  significant  estimated with the rule of mixture. Fig. 2 shows scanning  to note that  the surface oxidized specimen consists of  two  electron micrographs of  the polished and fracture surface  distinctly different  structural  scales, namely,  smooth and  of  the ZrB2-SiC composites. The microstructure of composites is regular, in which the mean grain size was about 6 lm and little porosity was observed in the polished  the  rugged oxides. The fine oxide grains with size of almost 4 lm are not closely adherent  to each other with respect  to the first oxide (Fig. 5b), while the latter contain many  section. The SiC particulate, which represents the majority  micro-cracks and voids. From Fig. 5c and d,  it  can be  of the dark features in Fig. 2,  is homogenously distributed  found that the latter may have been molten under the pres in the diboride matrix and no agglomeration was detected.  The growth of ZrB2 grains was restricted due to the existence of SiC particles.  perature  ent test conditions. It should be noted that the testing tem500 °C lower (2680 °C)  temperature of pure ZrO2  [13]. The mechanism  melting  almost  than  the  is  Fig. 2. SEM micrographs of the polished and fracture surface of ZrB2-20%SiC.  Fig. 3. ZrB2-20%SiC specimen before (left) and after (right) oxyacetylene torch testing at 2200 °C for 10 min.  \\x0c', 'for the formation of this molten scale is most likely due to  the solution of SiO2 in ZrO2 point of the ZrO2 (Fig. 6). The ported to the surface and then deposited at outside scale.  resulting in a lower melting  liquid phase was  trans The burst bubbles  and craterlets upon the  surface was  found (see Figs. 3 and 5), which presumably derives from  the evolution of volatile products  like SiO(g), CO(g) and  B2O3(g) during oxidation. The EDS analysis of the surface of a specimen oxidized at 2200 °C for 10 min shows the  presence of zirconium and oxygen as the primary elements,  while silicon and boron are not detected by the apparatus.  XRD examination of  the oxide surface (not  shown)  indi cates  that  the  surface  layer was primarily  composed of  monoclinic ZrO2.  The  highlight  of  the  cross-section  of  the  specimen  exposed to an oxidizing environment  is given in Fig. 7.  The analysis  revealed the presence of five distinct  layers,  in contrast  to the three and four distinct  layers  reported  for Hf-based and Zr-based materials [3,14], which may be  caused by the different oxidation temperatures  since  the  temperature in the present case was much higher than that  have been reported for the most oxidation studies. The outlayer was \\x1870 lm thick and composed of porous ZrO2, as has been observed in Fig. 7. Beneath this layer was found a dense recrystallized layer \\x1880 lm thick (layer 2) that has not been reported for Hf-based and Zr-based  ermost  materials in the previous studies [3,14]. Further analysis of  layer 2 (Fig. 8) showed that it consisted of large oxide particles with the size of approximately 20 lm. The formation  of this layer is thought to be the consequence of the high oxi dation temperature since high temperature is beneficial for  the  sintering of  zirconium oxide and the  growth of  the  grains. The third layer  is composed of many voids which  are probably formed due  to the active oxidation of SiC  and formation of B2O3 with a high partial pressure. The voids appear to be on the order of 25 lm which have not  been observed in lower temperature testing. This is a transi tional region between the recrystallized oxide and the under lying  SiC-depletion  region  (layer  4). Apparently,  the  bonding strength at  the  interface between this  layer and  the SiC-depletion layer is weakest among all  layers due to  the existence of numerous  large voids and the coefficient  of the thermal expansion mismatch between the ZrO2 (10.8 · 10\\x006 K\\x001) [15] and unaltered ZrB2 (6.7 · 10\\x006 K\\x001) [16]. For weak bonding and high vapor pressure, cracks  are easily formed and spallation is liable to occur. Thus, this  layer plays a very important role in the performance of the  0  100  200  300  400  500  600  1000  1200  1400  1600  1800  2000  2200  2400  T  e  m  p  e  r  a  t  u  r  e  ˚ (  C  )  Time (s)  Fig.  4. Measured  results  of  temperature  curves  of  ZrB2-20%SiC  composites.  Fig. 5. Surface morphologies of the oxide scale after ZrB2-20%SiC specimen exposure to oxyacetylene torch testing at 2200 °C for 10 min.  802  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806    \\x0c', 'J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  803  Fig. 6. Phase diagram of the binary ZrO2-SiO2 system [13].  material at high temperature and the control of its behavior  by optimizing the ratio of ZrB2 to SiC becomes essential to the improvement of the oxidation property at high tempertemperatures above 2200 °C. Under ature, especially for  neath this oxide  scale  is  the SiC-depletion region, which  Fig. 8. SEM micrograph of the second layer in Fig. 7.  was similar to the high temperature oxidation behavior of  the ZrB2-SiC and HfB2-SiC composites reported in the previous studies [3,14]. According to the EDS analysis, this  layer is very poor in silicon, which indicates that once the sil icon carbide is oxidized, the Si diffuses directly towards the  outside scale. The SEM and EDS results indicated that  it  dominantly consisted of unaltered ZrB2, whose shape and structure were not changed significantly during the oxida tion processing and the 3D interconnected nature remained.  The fifth layer is the unaltered ZrB2-SiC. Oxidation of ZrB2-SiC at 2200 °C in the oxyacetylene torch testing suggests that ZrO2 becomes a protective layer  Fig. 7. SEM result for a cross-section of ZrB2-SiC after oxidation at 2200 °C for 10 min.  \\x0c', 'at these temperatures by recrystallizing into a dense coher ent  scale. The variation of  the  thickness of  the  reaction  layer with time can be an indicator of the oxidation mech anism. The parabolic trend (Fig. 9) suggests reaction diffu sion-controlled kinetics.  In this  case,  the dense  coherent  ZrO2 surface layer acted as a barrier for diffusion of oxygen to the underlying ZrB2-SiC.  3.3. Oxidation mechanism of ZrB2-SiC  There are three kinds of oxidants,  i.e., H2O, CO2 and case. H2O and CO2 were not taken into the sake of simplicity. Apparently,  O2, account  in present  for  the  exis tence of H2O will the SiO2 and B2O3 (i.e., HBO2, H3BO3, Si(OH)4, SiO(OH)2 and SiO(OH)) leading to increased oxidation rate. Therefore,  increase  the volatilization of  the main expected reactions during oxidation process are  as follows  ZrB2 ðsÞ þ 5 O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðlÞ B2O3 ðlÞ ! B2O3 ðgÞ 2  ð3Þ  ð4Þ  SiCðsÞ þ 3 O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ 2 SiO2 ðlÞ ! SiO2 ðgÞ SiO2 ðlÞ ! SiOðgÞ þ 1=2O2 ðgÞ  ð5Þ  ð6Þ ð7Þ ð8Þ  Oxidation resistance depends on a protective oxide scale  formed on the surface of  the matrixes. Whether the oxide  products can form a dense adherent scale is very important  for ultra-high-temperature applications. In some cases, this  depends on the gas pressure in the internal oxide scale. The  vapor pressures of  the predominant gases  for ZrB2-SiC composites were calculated to compile a volatility diagram at 2200 °C (Fig. 10).  In air, B2O3(g) vapor species and its pressure is 2.7 atm (reaction 4). How is  the predominant  ever, at low oxygen partial pressure, SiO(g) has the highest  vapor  pressure  up to  12.3 atm (reaction  6). They  both  exceed 1 atm.  If  the formed gaseous phase products with  high partial pressures  cannot  escape  rapidly through the  outside scale and the gas pressure exceeds ambient prescreated. At 2200 °C,  sure, voids would be  the oxidation  resistance of SiC is poor due  to active oxidation, which  results in the formation of high vapor pressure oxide prod ucts like SiO and CO with almost no silica, which provides  effective oxidation protection at elevated temperature. The  active oxidation of SiC is a key issue in the use of SiC con taining composites at extremely high temperatures.  The  thermodynamic  stability  diagram for ZrB2-SiC temperature range of 1500-  with different 2500 °C was  P O2  in  the  calculated (Fig.  11), designating  the  stable  phases in each region to interpret and analyze the oxidation  mechanism of this kind of ultra-high-temperature ceramics. At any temperature, ZrB2 and SiC are stable at P O2 values below both lines. In the range of investigated temperature, the equilibrium P O2 for ZrB2 is higher than SiC, that ZrB2 is more stable than SiC at ultra-high-temperature. The selective removal of SiC from this region suggests  implying  -40  -30  -20  -10  0  10  -30  -20  -10  0  10  ZrB2+SiO2  BO2  B2O3(g)  SiO(g)  ZrO2+SiO2  ZrB2+SiC  g o L  P  a  r  t  i  a  l  p  r  s s e  u  r  e  Log pO 2  Fig. 10. Volatility diagram for ZrB2-SiC at 2200 °C.  10  20  30  40  200  300  400  500  600  700  800  900  Y =48.75+36.375 X-0.4125 X2  h T  i  k c  n  s e s s e  o  f  t  h  e  r  c a e  t  i  n o  l  e y a  r  (  µ  m  )  Holding time (min)   Fig. 9. Variation of the reaction layer thickness with time for exposure of ZrB2-SiC at 2200 °C.  -18 1800  1900  2000  2100  2200  2300  2400  2500  -16  -14  -12  -10  -8  -6  -4  ZrB2 , SiO and CO  ZrO2 , B2O3 , SiO and CO  ZrB2 and SiC  g o L  P  O  2  (  a  t  m  )  Temperature (°C)   SiC+3/2O2  ZrB2+5/2O2  SiO2+CO ZrO2+B2O3  → →  Fig. 11. Thermodynamic P O2 at different  stability diagram for ZrB2-SiC with different  temperatures.  804  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806                \\x0c', 'that  the active oxidation of SiC occurred without  signifi cant oxidation of ZrB2. provided by the thermodynamic calculations (Fig. 11). At 2200 °C, the equilibrium for \\x1810\\x0010 atm while the equilibrium P O2 for ZrB2 oxidation is \\x1810\\x009 atm. During oxidation, the P O2 in the depletion layer, including the interface between the depletion layer  Justification for  this assertion is  P O2  SiC  oxidation  is  and the unaltered ZrB2-SiC would be set by the oxidation of SiC thereby protecting ZrB2 from attack. The main reaction products from the oxidation of ZrB2- SiC composites are ZrO2, SiO2 and B2O3. The performance of the ZrB2-SiC ultra-high-temperature ceramic and the determination of whether the oxide products can form a  stable oxide scale are important for ultra-high-temperature  applications. This  is  largely dependent on volatilization  and decomposition of the oxide products. The vapor pres sure vs.  temperature above some oxides at ambient pres sure was  calculated  as  shown  in Fig. 12. The vapor 2200 °C are 2.8 · 105 Pa  pressures of B2O3 and SiO2 and 18 Pa, respectively, whereas  at  the  vapor  pressure  of  ZrO2 In fact, both SiO2 and B2O3 have high vapor pressures, so at high temperature they both quickly  is 0.025 Pa.  vaporize.  In addition, SiO2 could decompose by another mechanism, such as reaction 8. Furthermore, the formation  of some gaseous oxidation by-products (i.e., CO(g)), espe cially at high partial pressures, would accelerate the volatil ization of  the initial  formed glass and create short-circuit  paths for the incoming oxygen resulting in increased oxida tion rate. The volatilization of ZrO2 was negligible comparing with the other two oxides in this case since its vapor  pressure  is  very  low,  almost  three orders of magnitude  lower  than SiO2. The mass active oxidation of SiC and the volatilization of the partial  loss  closely correlates  to the  oxidation products such as B2O3 and SiO2. The protective role of the oxide products  is different,  which was mostly determined by the temperature. At low temperature (below 1100 °C), ZrB2 oxidizes much more rapidly than SiC and the formed liquid B2O3 acts as a barrier to  oxygen diffusion resulting in passive oxidation of ZrB2-SiC composites. Above 1300 °C, B2O3 becomes non-protective because of rapid evaporation and the higher temperature  protection was ascribed to SiO2. SiC significantly improves the oxidation resistance of the composites due to the forma tion of the more viscous silica-containing glass, which has a  higher melting temperature, a lower vapor pressure, and is a  greater a barrier to oxygen diffusion. Moreover, SiC signifi cantly increases the temperature range of the glass as a pri mary oxygen barrier, and extends  the maximum working  temperature in oxidizing environments. In fact, SiC-contain ing ZrB2 based UHTCs are effective up to high temperatures of 1700-1800 °C. Nonetheless, at higher temperatures, SiC is  no longer responsible for the improvement in oxidation resis tance of the composites because of a lack of silica formation  during the oxidation process. On one side, SiC undergoes a  transition from passive to active oxidation where the protec tive SiO2 layer is removed as SiO and CO. On the other side, the initially formed silica-enriched glass will be lost by con siderable evaporation and decomposition (reactions (7) and (8)). For temperatures higher than 1800 °C, ZrO2 recrystallizes into a dense coherent subscale (Figs. 7 and 8). This  subscale composed of zirconia appears to be able to protect  the underlying ceramic from catastrophic oxidation.  4. Conclusions  The ZrB2-20 vol%SiC composites exhibited excellent oxidation resistance at 2200 °C. The thickness of the reaction layer is only 375 lm after 10 min oxidation. The mass  and linear oxidation rates of the ZrB2-SiC composites are \\x000.23 mg/s and 0.66 lm/s, respectively. No macro-cracks or spallation were detected after oxidation. SiC exhibited  preferential oxidation compared with ZrB2 study. The formation of high pressure gaseous  in the present  phases  (i.e., B2O3 and SiO) amount of pores. SiC was no longer  led to the generation of a significant  responsible  for  the  improvement  in oxidation resistance of  the composites  in  the present case. ZrO2 recrystallized into a dense coherent subscale, which protected the underlying ceramic from cat astrophic oxidation.  Acknowledgements  This work was supported by the National Natural Sci ence Foundation of China (90505015 and 50602010),  the  Research Fund for the Doctoral Program of Higher Edu cation (20060213031) and the Program for New Century  Excellent Talents in University.  References  [1] Kuriakose AK, Margrave JL. The oxidation kinetics of  zirconium  diboride and zirconium carbide at high temperatures. J Electrochem  Soc 1964;111(7):827-31.  [2] Monteverde F. The thermal  stability in air of hot-pressed diboride  matrix  composites  for uses  at ultra-high-temperatures. Corros Sci  2005;47(8):2020-33.  1000  1200  1400  1600  1800  2000  2200  2400  2600  10 -10  10 7 10 6 10 5 10 4 10 3 10 2 10 1 10 0 10 -1 10 -2 10 -3 10 -4 10 -5 10 -6 10 -7 10 -8 10 -9  ZrO2(g)  SiO2(g)  B2O3(g)  V  a  o p  r  p  r  s s e  u  r  e  (  P  a  )  Temperature (˚C)  Fig. 12. Vapor pressure vs. temperature above some oxides, calculated at  ambient pressure.  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  805      \\x0c', '806  J. Han et al.  / Composites Science and Technology 68 (2008) 799-806  [3] Chamberlain AL, Fahrenholtz WG, Hilmas G, Ellerby D. Oxidation  [10] Opila EJ, Halbig MC. Oxidation of ZrB2-SiC. Ceram Eng Sci Proc  of ZrB2-SiC ceramics under atmospheric and re-entry conditions.  2001;22(3):221-8.  Refract Appl Tran 2005;1(2):1-8.  [11] Rezaie Alireza, Fahrenholtz WG, Hilmas GE. Evolution of structure  [4] Chamberlain AL,  Fahrenholtz WG, Hilmas GE,  Ellerby DT.  Characterization  of  zirconium diboride  for  thermal  protection  during the oxidation of zirconium diboride-silicon carbide in air up to 1500 °C. J Eur Ceram Soc 2007;27:2495-501.  systems. Key Eng Mater 2004;264-268(I):493-6.  [5] Opeka MM, Talmy IG, Zaykoski JA. Oxidation-based materials selection for 2000 °C + hypersonic aerosurfaces:  theoretical  consid [12] Paul Kolodziej. Aerothermal performance  constraints  for hyperve locity  small  radius  unswept  leading  edges  and  nosetips. NASA  Technical Memorandum 112204; 1997.  erations and historical experience. J Mater Sci 2004;39(19):5887-904.  [13] Butterman WC, Foster WR. Zircon stability  and the ZrO2-SiO2  [6] Chamberlain AL, Fahrenholtz WG, Hilmas GE, Ellerby DT. High phase diagram. The Am Mineral 1967;52:880-5.  strength  zirconium diboride-based  ceramics.  J Am Ceram Soc  [14] Fahrenholtz WG. Thermodynamic analysis of ZrB2-SiC oxidation:  2004;87(6):1170-2.  formation  of  a  SiC-depletion  region.  J  Am  Ceram  Soc  [7] Monteverde F, Bellosi A. Oxidation of ZrB2-based ceramics  in dry  2007;90(1):143-8.  air. J Electrochem Soc 2003;150(11):B552-9.  [15] Li XT, Shi JL, Zhang GB, Zhang H, Guo QG, Liu L. Effect of ZrB2  [8] Van Wie DM, Drewry  Jr DG, King DE, Hudson CM.  The  on  the  ablation  properties  of  carbon  composites. Mater  Lett  hypersonic  environment:  required operating conditions and design  2006;60:892-6.  challenges. J Mater Sci 2004;39(19):5915-24.  [16] Loehman R, Corral E, Dumm HP, Kotula P, Tandon R. Ultra-high [9] Monteverde F, Bellosi A. Microstructure and properties of an HfB2-  temperature  ceramics  for  hypersonic  vehicle  applications.  Sandia  SiC composite  for ultrahigh-temperature  applications. Adv Eng  Report. 2006;1-46.  Mater 2004;6(5):331-6.  \\x0c']"
},{
  "_id": 204,
  "PDF": "Oxyacetylene torch testing and microstructural characterization.pdf",
  "Text": "['Journal of Microscopy, Vol. 250, Pt 2 2013, pp. 122-129  Received 21 November 2012; accepted 15 February 2013  doi: 10.1111/jmi.12028  Oxyacetylene torch testing and microstructural characterization of tantalum carbide  A . P A U L ∗ ,  J . G . P . B I N N E R ∗ , B . V A I D H Y A N A T H A N ∗ , A . C . J . H E A T O N † & P . M . B R O W N †  ∗Department of Materials, Loughborough University, Loughborough, Leicestershire, UK †DSTL, Porton Down, Salisbury, UK  Key words. Oxyacetylene, tantalum carbide, ultra-high temperature.  Summary  Tantalum carbide samples have been subjected to hightemperature testing at 2300 C using an oxyacetylene torch to evaluate their potential for ultra-high temperature applications. While large samples cracked during the rapid heating,  indicating their inability to withstand thermal shock, small samples survived the severe test conditions. The oxidation products formed were characterized and found to comprise different phases of Ta2O5 . The ultra-high temperature experienced by the samples resulted in the formation of many inter esting microstructures, including the formation of submicron sized grains, which has not been reported previously in the literature, as well as the expected evidence of melting.  Introduction  Tantalum carbide, TaC, is a refractory carbide with one of the highest known melting points (3985 C) and is considered as an ultrahigh temperature ceramic, UHTC (Andersson et al., 2000; Tulhoff, 2000). There are several subcarbides of the Ta-C system with lower carbon contents; however, the  monocarbide, TaC, is the most widely used phase. Its main application is in cemented carbide as its addition improves the high-temperature fatigue strength, thermal shock resistance,  hot hardness and resistance against cratering and oxidation (Andersson et al., 2000; Tulhoff, 2000). TaC is a potential candidate for ultra-high temperature applications such as leading  edges and nose caps for hypersonic vehicles, rocket thrusters, combustor linings, etc., because of its high melting point.  There are only limited studies on the high-temperature oxidation behaviour of TaC materials. Desmaison-Brut et al. (1997) compared the oxidation behaviour of hot isostatically pressed TaC and Ta2C under isothermal and isobaric conditions in a dynamic flow of pure oxygen using a microbalance at relatively low temperatures of up to 850 C. They reported  Correspondence to: Anish Paul, Department of Materials, Loughborough University, Loughborough, LE11 3TU, UK. Tel: +44-1509-223154; fax: +44-1509-223949;  e-mail: a.paul@lboro.ac.uk  superior oxidation resistance for Ta2C and the formation of an oxycarbide layer at the Ta2C-Ta2O5 interface, whereas no such oxycarbide was observed for TaC. Thermogravimetric analysis, at temperatures up to 1500 C, has been used to compare the oxidation behaviour of TaC, TaC+10 wt% TaB2 and pure TaB2 (Zhang et al., 2008). It was observed that the oxidation resistance of TaC was marginally improved by the addition of TaB2 . Sciti et al. (2009) studied the oxidation of TaC with separate additions of 15 vol% of MoSi2 and TaSi2 at 1600 C in static air, exposing the samples for 15 min. Both materials were found to undergo severe oxidation, which con siderably reduced the strength of the composites. There are no reports on the ultra-high temperature testing of TaC ceramics using an oxyacetylene flame and the result ing morphology of the microstructure. This paper compares the oxidation behaviour and microstructure at and around the ablation centre of TaC samples subjected to oxyacetylene torch testing at temperatures up to 2300 C.  Experimental methods 100 × 100 × 10 mm TaC blocks were obtained from Cerac Inc., Milwaukee, USA. According to the manufacturer, the sample was densified via hot pressing and had a purity of  99.5%. The density of the sample was measured using the Archimedes principle in water and was found to be 93% of theoretical. 15 and 32 mm diameter discs for oxyacetylene  torch testing were machined out of the ceramic block using electrical discharge machining (EDM). A purpose built oxyacetylene torch test rig was used to per form the high-temperature oxidation testing. A schematic of the torch test facility is shown in Figure 1. Oxygen and acetylene in a 1.35:1 ratio was fed through a welding nozzle (No.13  nozzle, BOC Industrial, Surrey, UK) to produce an oxidizing flame. For testing, the samples were positioned in a water  cooled graphite holder that was then manually moved slowly into the hot zone of the flame. The samples were held there for 60 s and then the flame was extinguished allowing the sam ples to cool down naturally. During the test, the temperature  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society  \\x0c', 'O X Y A C E T Y L E N E T O R C H T E S T I N G A N D M I C R O S T R U C T U R A L C H A R A C T E R I Z A T I O N O F T a C  1 2 3  modal with an average size of 2-3 μm. Porosity is clearly visible as the ceramic was only 93% dense. From the XRD  pattern it can be seen that all the peaks correspond to the face centred cubic structure. The time temperature profiles  the 15 mm diameter  for  discs held at different distances from the nozzle during torch testing are given in Figure 3. It can be seen that increasing distance between the sample and nozzle resulted in much lower  heating rates and peak front face temperature, as expected. The heating rate for the first 5 s decreased from 330 C s−1 to 325 C s−1 to 275 C s−1 , whereas the peak front face temperature decreased from 2300 C to 2190 C and then to 1990 C as the distance increased from 10 mm to 15 mm and then to 20 mm respectively. It is well known that the gas flow rate and  gas flow ratio has a pronounced influence on the nature of the flame and its temperature; these two parameters were kept constant and only the distance was varied. The oxyacetylene  flame has an inner cone and an outer feather and the peak temperature is at the tip of the inner cone (Campbell, 1919). Figure 4 shows a view of one of the samples at 15 mm from  the tip of the nozzle, through a welding goggle, during testing. It can be seen that the inner cone of the flame is just touching  the sample surface. The length of this cone was determined to be 16 mm so this sample was exposed to the highest temperature. As well as a decrease in temperature, a decrease in gas  flow velocity was also expected with an increase in distance between the sample and the nozzle tip, though the magnitude of this could not be quantified.  Images of the samples after testing are shown in Figure 5. The front faces can be seen to be covered with a white layer indicating that the expected oxidation had occurred. The sur face oxide layer completely detached from the base material for the sample tested at 20 mm from the nozzle tip, whereas there was some adhesion for the other samples, possibly due  the formation of liquid phases at the test temperature and their subsequent solidification on cooling. However, no signs of surface erosion or cracking were observed for the 20 mm  sample, whilst erosion, melting and cracking were all observed for the other two distances. On cross-sectional examination,  the cracks were found to propagate into the bulk. The response of the 32 mm diameter TaC discs to the flame was completely different due to the large thermal gradients  that developed. Unlike the 15 mm diameter discs, which were embedded in a carbon-carbon holder, these samples were held directly by the graphite sample holder using three  graphite bolts. During the test, there was rapid heating and severe thermal gradients. As the samples tried to expand, the  graphite bolts resulted in stresses being generated that resulted  in crack formation. Three out of four 32 mm TaC discs tested shattered and fell out of the sample holder before reaching even 1000 C. While the fourth sample cracked, out allowing the 60 s test to be completed, Figure 6. The cracks originated from the centre of the sample, where the oxyacetylene flame was focused, and propagated to the points where  it didn’t fall  Fig. 1. Oxyacetylene torch test set up. (1) Back face thermocouple, (2)  water cooling, (3) sample holder, (4) sample, (5) guide rail, (6) protective  insulation, (7) oxyacetylene torch, (8) neutral density filter, (9) thermal  imaging camera and (10) two-colour pyrometer.  distribution was recorded using a modified infrared thermal imaging camera (Thermovision A40, FLIR Systems AB, Danderyd, Sweden) and the peak temperature was recorded us ing a Marathon MR1SCSF two-colour infrared pyrometer (Raytek GmbH, Berlin, Germany). The back face temperature  of the samples was also recorded during testing using a K-type thermocouple connected to a data logger. For testing the 15 mm diameter discs,  the samples were  placed in a carbon-carbon insert prior to fixing inside the graphite sample holder whereas the 32 mm samples were fixed directly within the holder. For the 15 mm diameter  discs, distances of 10, 15 and 20 mm were used between the  sample and the nozzle to alter the peak temperature and heating profile. Four 32 mm diameter samples were also tested at  10 mm distance, but three of them shattered within a few seconds of moving into the flame, even before reaching the peak temperature.  The surface microstructure of the TaC before and after oxyacetylene torch testing was characterized using a field emission gun scanning electron microscope (FEGSEM, LEO  1530VP, Carl Zeiss SMT, Oberkochen, Germany) and the chemical composition was examined using energy dispersive spectroscopy, EDS (EDAX, EDAX Inc., NJ, USA). The mi crostructure and composition beneath the oxide layer was also analysed after removing the oxide; the adhesion between the  base layer and oxide was poor and the latter was easily removed. X-ray diffraction patterns of the as-received TaC and the various oxidation products formed after the oxyacetylene  torch tests were recorded with a Bruker D8 model diffractometer (Bruker AXS GmbH, Karlsruhe, Germany) using Cu-Kα radiation. The samples were cross sectioned after the oxyacety lene test, polished and characterized using FEGSEM and EDS.  Results  The microstructure, EDS spectrum and XRD pattern of the TaC before the oxyacetylene torch test are shown in  Figure 2. It can be seen that the grain size distribution is uni C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', '1 2 4  A . P A U L E T A L .  Fig. 2.  (a) SEM, (b) EDS and (c) XRD of TaC before torch testing.  it was held by the bolts; highlighted by the arrows. One way to avoid this cracking in the future is to eliminate any stresses  from the sample holder. The thermal gradients on the 32 mm disc are shown in Figure 7 and the peak back face temperature was recorded to be 1050 C.  Fig. 3. Time-temperature profiles of TaC discs subjected to oxyacetylene  flame at 10, 15 and 20 mm from the nozzle.  The formation of a white layer was also observed on the 32 mm diameter sample front surface, together with straw  coloured, frozen liquid phases with a clear melt boundary. Xray diffraction patterns of both the white surface layer and the solidified molten structure are shown in Figure 8 (a) and the layer comprised α and β phases of Ta2O5 , with the strongest peak corresponding to that of the α . Similarly, the frozen droplets were also found to comprise mainly α , with  (b);  small amounts of β also being detected. Cross-sectional analysis of the 15 mm discs was performed  to compare the thickness of the oxide layers formed and the chemical compositions, Figure 9. An increase in oxide layer thickness was expected with an increase in temperature, but  a decrease was actually observed at the highest temperatures due to the melting of the Ta2O5 and its removal by the gas flow. This clearly highlights the disadvantage of using this type of  material at ultra-high temperatures when a high-velocity gas flow is present. EDS analysis of the frozen droplets and the oxide layer indicated the presence of Ta and O, whereas the  bulk, unaffected material continued to reveal only Ta and C. The surface of the disc after oxyacetylene torch testing can be divided into three separate zones, each with a different  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', 'O X Y A C E T Y L E N E T O R C H T E S T I N G A N D M I C R O S T R U C T U R A L C H A R A C T E R I Z A T I O N O F T a C  1 2 5  Fig. 4. View of sample through a welding goggle during oxyacetylene  torch testing; the inner cone of the flame is clearly visible. The sample was  15 mm from the nozzle tip.  Fig. 5. Photographic images of  the TaC discs after oxyacetylene torch  testing at (a) 20 mm, (b) 15 mm and (c) 10 mm distance from the nozzle  tip.  morphology, Figure 6. Zone 1 is at the flame tip, where the flame directly impinged the sample and hence the temperature was highest. Zone 2 is the flow zone; the temperature  experienced was lower and the molten products formed in zone 1 were swept across by the high-velocity gas flow from the torch nozzle. Zone 3 experienced the lowest temperature  and no melting was observed. A clear melt boundary is visible between zones 2 and 3. The detailed microstructural analysis of  the sample after  testing revealed several interesting features. The white oxide layer formed on the surface of the sample in zone 1 was found to be highly porous, Figure 10 (a). This is believed to be due to the escape of the CO/CO2 gases formed at high temperatures and the porosity allows easy diffusion of oxygen towards the bulk of the sample. The microstructure beneath the white oxide layer  in zone 1 is shown in Figure 10 (b). This region is not as porous as the surface layer, indicating only a partial oxidation and the individual grains were observed to have small pores within  them. EDS analysis of this region detected mainly Ta and O. If the sample was held at the testing temperature for longer, it is  believed that these grains would also have developed porosity similar to that of the surface layer. Many interesting microstructures were observed in zone  Fig. 6. 32 mm diameter TaC sample after an oxyacetylene torch test. The  graphite sample holder is also seen and the arrows indicate the points  where the sample was gripped using the bolts.  Inset picture shows the  TaC disc before testing. The numbers indicate the three different zones  observed after torch testing. (1) Flame tip zone, (2) flow zone and (3) outer  zone.  Figure 11(a) shows the formation of  large columnar grains  growing in a radial direction. These will have been formed as a result of the high-velocity flame pushing the molten oxide out of zone 1. A higher magnification view of these columnar  grains indicates the presence of very fine, dense looking grains, Figure 11(b), (c) and (d). EDS analysis (not shown) detected  only Ta and O in these fine grains. Near to the melt boundary, but within zone 2, the columnar grains were oriented in random directions because of the  rapid cooling, as shown in Figure 12. The segregation of a secondary phase was identified on some of the frozen droplets and EDS analysis confirmed the presence of Fe, which is be lieved to have come from the impurities present in the starting material, Figure 13, although no Fe was detected on the EDS of the starting material, Figure 2(b). It is assumed that the very  low quantities of Fe present in the starting material were uniformly distributed and during ultra-high-temperature testing they melted and coalesced in the Ta2O5 droplets. A similar type of segregation was reported by Sautereau & Mocellin (1974) after the thermal etching of sintered TaC and NbC samples at 1400 C. They reported that the impurities present in the starting powders coalesced and formed droplets, predominantly at the grain centres, during the thermal etching process. The use of Fe as a sintering aid during the hot pressing of TaC has also  been reported (Scholz, 1963).  Discussion  Oxidation of tantalum carbide has been observed to proceed with the formation of a porous, nonprotective oxide scale.  2,  the area between the flame tip and the melt boundary.  Melting of the scale is observed when the test temperature is  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', '1 2 6  A . P A U L E T A L .  Fig. 7. Thermograph showing the thermal gradients developed during the oxyacetylene torch testing of a 32 mm diameter TaC disc at 10 mm from the  nozzle. (a) Before cracking, (b) immediately after cracking and (c) before extinguishing the flame.  >1900 C. The major reactions taking place during the oxidation of TaC are given by equations 1 and 2 (Desmaison-Brut et al., 1997).  4TaC + 7O 2 → 2Ta2 O 5 + 4CO  2TaC + 9/2O 2 → T a 2 O 5 + 2CO2  (1)  (2)  The formation of other oxides, such as TaO2 and Ta2O3 , has also been reported at intermediate temperatures of around  Fig. 8. X-ray diffraction patterns on the 32 mm diameter sample after  oxyacetylene torch testing.  (a) White layer formed on the surface and  (b) solidified molten structure.  900-1300 C, depending on the oxygen partial pressure (Chen et al., 2010). No such oxides were observed after oxyacetylene torch testing, though they might have been formed during the  heating stages. The indexing of the X-ray diffraction pattern of the lowtemperature form of Ta2O5 has always confused crystallographers. Most of the intense diffraction lines can be indexed using an orthorhombic unit cell; however, many of the weak lines have not been interpreted unambiguously. The rea son for this is that the position of the weak lines is dependent on the nature of the heat treatment and the amount of cation and anion impurities (Roth et al., 1970). The sit uation is made more complex in this study because of the rapid heating and cooling involved. It has also been reported that the high-temperature polymorph (α -Ta2O5 ) undergoes several phase transformations on quenching. Although it is triclinic at room temperature, it is monoclinic at intermedi ate temperatures and tetragonal at the temperature at which it is stable (Waring & Roth, 1968). XRD patterns of the oxide layers from this study indicated the presence of both α and β -Ta2O5 . The crystal structure of the α -phase was observed to be triclinic. β -Ta2O5 is a low-temperature phase with a transition temperature of 1320 ± 20 C to the α phase, which is stable at high temperature (Lagergren & Magn ´eli, 1952). Reisman et al. (1956) reported a slightly higher transformation temperature of 1360 ± 5 C after heating Ta2O5 using an oxy-gas flame and also suggested that this transformation is sluggish but completely reversible. Quenching of the molten phases resulted in a straw coloured, crystalline Ta2O5 that had a similar structure to the frozen droplets from this study. For Reisman et al. (1956) however, their X-ray diffraction pattern of the solidified struc ture revealed the presence of just phase was also observed in this  the α -phase whereas β study. The presence of  the former is expected to be due to the rapid heating and cooling and the sluggish transformation behaviour. An increase in the α -phase intensity for the frozen droplets is  possibly due to the higher temperature experienced by the droplets and the minimal time available for this phase to revert to β due to rapid cooling. The presence of ε and δ  phases (Izumi & Kodama., 1979), and many high-pressure  phases (Filonenko & Zibrov, 2001; Zibrov et al., 2000), is also  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', 'O X Y A C E T Y L E N E T O R C H T E S T I N G A N D M I C R O S T R U C T U R A L C H A R A C T E R I Z A T I O N O F T a C  1 2 7  Fig. 9. Cross-sectional images from directly below the flame tip from the 15 mm diameter samples subjected to oxyacetylene torch testing at three different  distances from the nozzle, (a) 20 mm (b) 15 mm and (c) 10 mm. The EDS patterns in (d) were recorded from locations 1, 2 and 3 in Fig. 9b.  reported in the literature, but was not observed in this work. Even though the melting point of TaC is around 3985 C, it is reported to undergo oxidation at temperatures as low as 800 C (Zhang et al., 2008) and the resulting oxides melt  at around 1800 C (1785 ± 30 C for β -Ta2O5 and 1872 C for α -Ta2O5 (Reisman et al., 1956)), which is well below the oxyacetylene torch test temperatures used in this study.  Fig. 10. FEGSEM image of (a) white surface oxide layer, (b) partially oxidized grains below the surface layer, (c) the EDS pattern taken on the partially  oxidized grains in (b).  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', '1 2 8  A . P A U L E T A L .  Fig. 11.  (a) SEM image  showing the  formation of  columnar grains.  (b) Higher magnification view of the area highlighted by the rectangle in  (a). (c) and (d) are higher magnification images of the submicron grains.  Fig. 12. SEM images showing the random orientation of grains at  the  melt boundary.  The XRD results are not completely in agreement with those reported after the oxidation of C-TaC-C composites at temperatures up to 1400 C (Chen et al., 2010), where the formation of a hexagonal structure (δ -Ta2O5 ) was reported at temperatures below 900 C. This then transformed to an orthorhombic (β -Ta2O5 ) structure above this temperature. This difference  could be due to the large difference in testing temperatures and the more oxidizing conditions. Gaseous CO/CO2 is produced as a by-product during the oxidation of TaC. The escape of these gases produces continuous porosity in the oxide layer, as reported by Courtright et al.  (1991). The pores act as pathways for both the inward transport of oxygen and further outward diffusion of CO, preventing CO build up and catastrophic delamination of the oxide layer  under static oxidation conditions. The delamination of the oxide layer observed in this study was again partially due to the rapid heating rates involved, these will have limited the time  available for the gasses formed to escape. Further factors are outlined below. The pores present in the grains beneath the fully oxidized  surface layer will have been created by the partial internal oxidation of the TaC grains. Similar observations were made by Courtright al. (1991). A well-defined carbide-oxide  et  boundary layer indicated the reaction front. The oxidation reaction taking place at this interface can be assumed to control the rate of oxidation. Similar interfacial oxidation reactions  were also reported for other high-temperature carbides such as HfC and ZrC with an oxygen gradient from the oxide layer  to the carbide. The diffusion of oxygen into the carbide limits the rate of the reaction (Shimada, 2001). A change in the rate controlling mechanism at 1800 C has also been reported in the literature. The lower temperature kinetics can be explained by assuming that oxide growth is controlled by O2 , CO and CO2 gaseous transport through the interconnected pores in the oxide. This process became less important at increasing temperatures as the oxide sintered and the effective volume available for gaseous diffusion was reduced (Cour tright et al., 1991). During the oxyacetylene torch testing, the oxides formed melted making it more difficult for the O2 to diffuse through the viscous liquid phase compared to diffusing  through the interconnected porous oxide. Also, near the tip of the flame, the melt formed was pushed away from the flame  Fig. 13. FEGSEM image showing the formation of a secondary phase on the solidified droplets. EDS detected the presence of Fe along with Ta, O and C.  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  \\x0c', 'O X Y A C E T Y L E N E T O R C H T E S T I N G A N D M I C R O S T R U C T U R A L C H A R A C T E R I Z A T I O N O F T a C  1 2 9  front by the high-velocity gas flow hitting the sample, exposing fresh surfaces for oxidation and influencing the rate of the  reaction. The formation and escape of any intermediate oxides such as TaO2 and Ta2O3 will also have led to the generation of more porosity, helping the rapid inward diffusion of oxygen  and accelerating the oxidation process (Chen et al., 2010). In addition to gas build up, the difference in coefficient of thermal expansion between the oxide and the underlying carbide can also play a role in delamination. The CTE of Ta2O5 is reported 2.9 × 10−6  C−1 to be around from room temperature to 500 C (Moldovan et al., 2004) whereas 6.29 × 10−6  C−1 at room temperature (Tulhoff, 2000). that of TaC is The molten oxides recrystallized on cooling forming columnar grains. The formation of the submicron grains is expected  to be due to the formation of tantalum suboxides, TaO2 and Ta2O3 at intermediate temperature of oxidation, probably during the heating up stage. The formation temperature of some of these suboxides are reported to be <600 C (Chang & Phillips, 1969). It is believed that some of these oxides will have vaporized at the ultra-high temperatures being  including  used (Chen et al., 2010) and then deposited at regions of relatively lower temperature on the sample surface. They would  subsequently undergo sintering, leading to the formation of dense-looking submicron grains on the surface. The presence of molten phases is expected to act as nucleation sites for the  growth of these grains.  Conclusions  Even though TaC has the highest melting point amongst all  the potential materials for UHTC applications, the results of the oxyacetylene torch testing proved its inability to fulfil the requirements for such applications. TaC was unable to withstand  the high thermal shocks experienced during the test, leading  to severe cracking. It oxidized at temperatures well below the melting point of the carbide and the oxides melted at temperatures below 1900 C rendering these materials unsuitable for applications requiring ultra-high temperature (>2000 C) resistance. Moreover, the oxide layer formed was porous and  not completely adherent, offering little or no additional protection at elevated temperatures. The oxides were found to comprise the lowand high-temperature stable phases of Ta2O5 . The formation of submicron grains was expected to be due to the volatilization, deposition and subsequent sintering of the  oxide phases formed at intermediate temperatures.  Acknowledgements  The authors thank the UK’s Defence Science and Technology Laboratory (DSTL) for providing the financial support for this  work under contract number DSTLX-1000015267 as well as the US Air Force Research Laboratory’s Materials and Manu facturing Directorate for on-going collaborations.  C(cid:5) 2013 The Authors Journal of Microscopy C(cid:5) 2013 Royal Microscopical Society, 250, 122-129  References  Andersson, K., Reichert, K. & Wolf, R. (2000) Tantalum and tantalum  compounds. Ullmann’s Encyclopedia of Industrial Chemistry, pp. 1-15.  Wiley-VCH Verlag GmbH & Co. KGaA, New York.  Campbell, L. (1919) Chapter II: operation. Oxy-Acetylene Welding Manual,  pp. 27-38. John Wiley and Sons, Inc., New York.  Chang, L.L.Y. & Phillips B.  (1969) Phase relations in refractory metal oxygen systems. J. Am. Ceram. Soc. 52, 527-533.  Chen, Z., Xiong X., Li G. & Wang Y. (2010) Mechanical properties and  oxidation behaviors of carbon/carbon composites with C-TaC-C multi interlayer. J. Mater. Sci. 45, 3477-3482.  Courtright E.L., Prater  J.T., Holcomb G.R., Pierre G.R.S. & Rapp R.A.  (1991) Oxidation of hafnium carbide and hafnium carbide with addi tions of tantalum and praseodymium. Oxid. Met. 36, 423-437.  Desmaison-Brut M., Alexandre N. & Desmaison J. (1997) Comparison of  the oxidation behaviour of two dense hot isostatically pressed tanta lum carbide (TaC and Ta2 C) materials. J. Eur. Ceram. Soc. 17, 1325-  1334.  Filonenko, V.P. & Zibrov I.P. (2001) High-pressure phase transitions of M2O5 (M = V, Nb, Ta) and thermal stability of new polymorphs. Inorg. Mat. 37, 953-959.  Izumi, F. & Kodama H. (1979) A new modification of tantalum (V) oxide.  J. Less Comm. Metals 63, 305-307.  Lagergren, S. & Magn ´eli A. (1952) On the tantalum-oxygen system. Acta.  Chem. Scand. 6, 444-446.  Moldovan, M., Weyant C., Johnson D. & Faber K. (2004) Tantalum oxide  coatings as candidate environmental barriers. J. Therm. Spray Technol.  13, 51-56.  Reisman, A., Holtzberg F., Berkenblit M. & Berry M. (1956) Reactions of  the group VB Pentoxides with alkali oxides and carbonates. III: thermal  and X-ray Phase diagrams of the system K2O or K2 CO3 with Ta2O5 . J. Am. Chem. Soc. 78, 4514-4520.  Roth, R.S., Waring J.L. & Parker H.S. (1970) Effect of oxide additions on  the polymorphism of  tantalum pentoxide.  IV. The  system Ta2O5  Ta2WO8 . J. Solid State Chem. 2, 445-461.  Sautereau, J. & Mocellin A. (1974) Sintering behaviour of ultrafine NbC  and TaC powders. J. Mater. Sci. 9, 761-771.  Scholz, S. (1963) Some new aspects of hot pressing of refractories. Special Ce ramics 1962, Proceedings of a Symposium Held by the British Ceramic  Research Association, 293-307.  Sciti, D., Silvestroni L., Guicciardi S., Fabbriche D.D. & Bellosi A. (2009)  Processing, mechanical properties and oxidation behavior of TaC and  HfC composites containing 15 vol% TaSi2 or MoSi2 . J. Mater. Sci. 24, 2056-2065.  Shimada, S.  (2001)  Interfacial  reaction on oxidation of carbides with  formation of carbon. Solid State Ionics 141-142, 99-104.  Tulhoff,  H.  (2000)  Carbides.  Ullmann’s  Encyclopedia  of  Industrial  Chemistry,  pp. 555-582. Wiley-VCH Verlag GmbH & Co. KGaA,  New York.  Waring, J.L. & Roth R.S. (1968) Effect of oxide additions on the polymor phism of tantalum pentoxide (System Ta2O5 -TiO2 ). J. Res. Nat. Bur. Stand. Sect. A. 72A, 175-186.  Zhang, X., Hilmas G.E. & Fahrenholtz W.G.  (2008) Densification, me chanical properties, and oxidation resistance of TaC-TaB2 ceramics. J. Am. Ceram. Soc. 91, 4129-4132.  Zibrov, I.P., Filonenko V.P., Sundberg, M. & Werner P.-E. (2000) Struc tures and phase  transitions of B-Ta2O5 and Z--Ta2O5 : pressure forms of Ta2O5 . Acta Crystallog. Sec. B 56, 659-665.  two high \\x0c']"
},{
  "_id": 205,
  "PDF": "Oxygen diffusion mechanisms during high-temperature oxidation of ZrB2-SiC.pdf",
  "Text": "['O R I G I N A L A R T I C L E  Oxygen diffusion mechanisms during high-temperature oxidation of ZrB2-SiC  Kathleen S. Cissel1,2  |  Elizabeth Opila2  1UES,  Inc., Dayton, Ohio  2University of Virginia, Department of  Materials Science and Engineering,  Charlottesville, Virginia  Correspondence  Kathleen S. Cissel, UES,  Inc., Dayton,  OH.  Email: kns9a@virginia.edu  Funding information  National Hypersonic Science Center -  Materials and Structures  Abstract  Oxygen diffusion mechanisms during oxidation of ZrB2-30 vol% SiC were explored at temperatures of 1500°C and 1650°C using an 18O tracer technique. Double oxida tion experiments in 16O2 and 18O2 were performed using a modified resistive heat ing  system. A combination  of  scanning  electron microscopy,  energy-dispersive  spectroscopy, and time-of-flight secondary ion mass spectrometry was used to char acterize the borosilicate and ZrO2 oxidation products. Oxygen exchange with the  borosilicate network was observed to occur quickly at  the oxygen-borosilicate sur face at both 1500°C and 1650°C, while evidence of oxygen permeation was only  observed at 1650°C for short  time (<1 min) exposures. At  longer  times, >5-9 min,  complete oxygen exchange  throughout both the borosilicate glass and ZrO2 was  observed at both temperatures preventing identification of  the oxygen transport  mechanisms, but demonstrating that oxygen transport is rapid in both oxide phases.  K E Y W O R D S  diffusion/diffusivity, oxidation, ultrahigh-temperature ceramics  1  |  I N T R O D U C T I O N  Ultrahigh-temperature ceramics (UHTCs) are a family of materials that have melting temperatures in excess of 3000°C,  making them attractive for aerospace applications such as ther mal protection systems  (TPS) on hypersonic flight vehicles.  Design of future hypersonic vehicles will  incorporate sharper  wing  leading  edges  and  nose cones, allowing speeds.1,2 As  for  better  maneuverability and higher  the  leading edge  radius of a component  is made sharper and the speed of  the  vehicle is increased,  the leading shockwave will  tend toward  the surface. This close proximity of the shock and increase in  near surface velocity gradient greatly increase the stagnation  point heat  flux resulting in operational  temperatures of these 1700-3000°C.2-4 This  components  in  the  neighborhood  of  temperature range is beyond the operating capability of cur rent TPS, necessitating the development of new TPS materials.3-8 One of the largest concerns for nonoxide TPS materials  is oxidation behavior at ultrahigh temperatures. Studies on the  oxidation characteristics of the UHTC ZrB2-SiC have been conducted since the mid-1960s,3,4,8-10 however, the oxidation  mechanisms are still not completely understood.  ZrB2-SiC forms a complex oxide morphology, which is dependent on oxidation temperature.11 The oxidation reac tions are:  ZrB2  sð  Þ þ 5  2  O2 gð  Þ ¼ ZrO2  sð  Þ þ B2O3 ðl or gÞ  (1)  SiC sð  Þ þ 3  2  O2 gð Þ ¼ SiO2 ðsÞ þ COðgÞ Þ ¼ SiO(g) þ CO(g) Þ ¼ SiO2 ðsÞ þ C(s) lð Þ ¼ xSiO2 \\x01 yB2O3  (2a)  SiC sð SiC sð xSiO2 ðs,lÞ þ yB2O3  Þ þ O2 gð Þ þ O2 gð  (2b)  (2c)  s,  l  ð  Þ  (3)  SiO gð  Þ þ 1  2  O2 gð  Þ ¼ SiO2 ðsÞ  (4)  CO gð  Þ þ 1  2  O2 gð  Þ ¼ CO2 ðgÞ  (5)  Kathleen Shugart Cissel and Elizabeth Opila are American Ceramic Soci ety members.  Received: 17 April 2017  |  Accepted: 16 October 2017  DOI: 10.1111/jace.15298  J Am Ceram Soc. 2018;101:1765-1779.  wileyonlinelibrary.com/journal/jace  © 2017 The American Ceramic Society  |  1765  \\x0c', 'Prior work by these authors11 has shown that at temperatures below 1627°C, a two-layer oxide forms which conZrO2+C beneath a temperatures of 1627°C and  sists  of  borosilicate  glass  layer. At  above,  a borosilicate glass a ZrO2+borosilicate layer of ZrB2 depleted of SiC is formed below the oxides. Five possible pathways/mechanisms  layer  is  found  at  the  surface  above  layer. A porous  for  oxygen  diffusion  in  oxides  grown  on  ZrB2-SiC can envisioned due to the multiple phases and morphologies in  be  the oxide,  illustrated in Table 1.  It has been theorized by many (including Parthasarathy et al)12,13 that  the rate controlling mechanism of  the oxida tion of ZrB2-SiC is the inward diffusion of oxygen through the borosilicate glass, while the ZrO2 grains have negligible contribution, however, this has not been decidedly proven.9,14-18 Molecular oxygen may permeate through intersti tial spaces in the amorphous borosilicate structure. Network  exchange,  similar  to lattice diffusion in crystalline phases,  may occur, whereby the oxygen jumps from one network site  to the next in the borosilicate glass. Deal and Grove show permeation of oxygen in thermally grown oxide on silicon.19  Ogbuji and Opila have found that  the enthalpy for the oxida tion of SiC is 1500°C,  essentially  the  same  as  that  for Si  up  to  indicating  that  the  controlling  mechanism  of  oxidation is the same for both materials.20 Network diffusion  is shown to be much slower  than permeation during oxida tion  of  silicon  and  silicon  carbide  as well  as  in  silicate  glasses, though the difference lessens as temperature increases.21-23 Zheng et al have found a transition from permeation to exchange to occur above 1200°C for single-crystal SiC.24,25 For vitreous silica, Kalen et al  found that both operation.26  network  and  interstitial mechanisms were  in  However,  the authors are unaware of any studies that have  been conducted at temperatures as high as those studied here (1500-1650°C), nor of any on borosilicate glasses.  Much of  the research for oxygen diffusion in ZrO2 has been conducted on Yor Ca-doped ZrO2 27 and monoclinic ZrO2, which is stable at room temperatures, while at high temperatures such as those studied here, stable phase.18,28 The  tetragonal ZrO2 addition of Y or Ca dopants  is  the  is  known to increase oxygen diffusion through vacancy formation.27,29  In addition, evidence of grain boundary diffu sion was found by Brossmann et al undoped monoclinic ZrO2.28-30 In tetragonal ZrO2, oxygen is believed to diffuse by the vacancy mechanism, due to intrinsic oxygen vacancies.18,31 There  for ultrafine grained,  the predominance of  may also be Knudsen diffusion diameter pores in the material.32,33  occurring  through  small  T A B L E 1  Possible oxygen diffusion pathways/mechanisms during oxidation of ZrB2-SiC  M  c e  h  a  n  i  s  m  Transport through borosilicate glass   Transport  through ZrO2 Lattice or  Grain  Boundary  Diffusion  Transport through  pores Knudsen Diffusion  Permeation  Network Exchange  A  t  o  m  i  s  t  i  c  s  o  f  D  i  f f  u  s  i  o  n  No breaking of bonds.  Molecular O2 is being  transported25.  Molecular O2 must  dissociate into 2O- 25  .  Molecular O2 must dissociate   into 2O--31, 32 .  Molecular O2 travels  through pores34, 35.  Lattice:  Grain Boundary:  ZrO2 grain  ZrO2 grain  Grain Boundary  1766  |  CISSEL AND OPILA      \\x0c', 'The objective of  this work is  to explore  the diffusion  pathways of oxygen during oxidation of ZrB2-30 vol% SiC below and above the onset temperature for SiC depletion (1500°C and 1650°C) using 18O tracer diffusion tech18O tracer diffusion through thermally grown 16O  niques.  oxide layers  (double oxidation)  is a powerful  technique to  identify oxygen diffusion pathways during an oxidation process.34-37 Double oxidation experiments oxidize a samin 16O2 and then in 18O2 to grow oxide using both isotopes. The diffusivity varies with the square root of the mass,38  ple first  therefore,  the  18O  and  16O  diffusivities  are  expected to vary by about 6%. Given the difficulty in mea suring oxygen diffusion coefficients  accurately,  as well  as  all  other  uncertainties  in modeling materials  responses  in  actual  flight  space,  it  is  concluded  that  this  difference  is  negligible. The information gained contributes to improved  understanding of oxidation mechanisms,  improved life pre diction,  and  potentially  identification  of  strategies  for  improving oxidation resistance of ZrB2-SiC.  2  |  E X P E R IM E N T A L P R O C E D U R E  ZrB2-30 vol% SiC specimens were University of Science and Technology using attrition milled powders which were then hot pressed.39,40 Specimens were  fabricated at Missouri  cut  from a series of billets using an automated surface grin der  into  bars  of  40 mm 9 4 mm 9 3 mm and were  fin ished  using  1200  grit  diamond  abrasive. The  bars were  shipped to Bomas Machine Specialties (Somerville, MA) for machining into “bridge”-shaped specimens for oxidation testing. The bridges were 15 9 3 9 1.5 mm, with the mid dle 3.5 mm thinned to 0.5 mm (Figure 1). Prior  to oxida tion, all  specimens were ultrasonically cleaned using baths  of  detergent  and DI water, DI water,  acetone,  and  then  ethanol.  Resistive heating was used to test the oxidation behavior at ultrahigh temperatures (1300-1800°C), after a method  first developed by Karlsdottir and Halloran, which they called “the ribbon method.”41 This method resistively heats the specimen using a high current, ~100 amps, that runs  through the specimen. As the resistance of ZrB2 ~15 9 106 Ωcm,42 this approach works well. specified specimen geometry  is  similar  to metals,  The  (Figure 1)  allows  the  center of  the specimen, which is not  in direct contact with  any surfaces,  to heat  to ultrahigh temperatures, while  the  supporting  ends  remain  relatively  cool, minimizing  high temperature contamination and interaction of  the specimen  with  the  heating  apparatus.  The  system developed  by  Karlsdottir  and Halloran was modified so that  the  sample  could  be  enclosed  in a small volume trolled atmosphere exposures.11 The chamber volume was optimized such that minimal 18O2 could be used while still providing sufficient amount of isotope for growth of new  chamber  for  con oxide. This method has  the advantage of being more eco nomical  than  high-enthalpy  plasma  facilities  (arc-jet  and  inductively coupled plasmas)  and allows  for more  control  than  using an oxyacetylene torch, reaching T > 1500°C.  both  alternative  tech niques for  For  testing, each end of  the bridge specimen was posi tioned on a section of  flattened copper wire with a cross section of 1.05 9 2.85 mm, which was attached to the cur rent  feedthroughs  in  the  specimen  chamber. A piece  of  platinum foil  (0.06 mm thick, Heraeus Materials Technol ogy, Chandler, AZ) was placed  between the  copper wire  and the specimen to minimize reactions between the speci men and the wire. The  specimen was held in place using  toothless  alligator  clips, which were  electrically  isolated  using 0.8-mm thick sheets of Maycor  (a glass-mica  cera mic, McMaster-Carr, Robbinsville, NJ) or  stabilized ZrO2 Inc., Houston, TX). A bar of ZrO2 was glued to the side of the specimen to provide stability  (Ceramic Technologies,  against  torque when inserting the specimen into the system.  B2O3(g) flowing O2(g) and a rate of 900 sccm. The O2(g) (CaSO4) prior to entering the resistive heating chamber maintain low humidity. The current was controlled using a  condensation  on  the window was  prevented  by  through the chamber at a pressure of 1 atm  flowed through drierite  to  Eurotherm controller  (BPAN controller, Micropyretics Hea ters  International,  Inc., Cincinnati, OH). The  current  run ning through the specimen was read continuously by a NI 6009 data acquisition card (DAQ). The specimen tempera ture was measured using an emissivity correcting infrared  pyrometer  (Pyrofiber  Lab  PFL-0865-0790-2500C311,  Pyrometer  Instrument Company, Windsor, NJ) with a spot  size  of  1.5 mm, which  automatically  corrected  for  the  changing emissivity (generally 0.80-0.99) of  the oxidizing  surface using pulsed laser  technology. The pyrometer was ~4 measurements  set  to  a  data  acquisition  rate  of  per  F I G U R E 1  ZrB2-30 vol% SiC specimen machined into bridge  shape for resistive heating with the middle 3.5 mm thinned to 0.5 mm  CISSEL AND OPILA  |  1767  \\x0c', 'second and wire melting point operates within 28°C of true temperature in the test (1500-1800°C),  tests indicate that  the system  temper ature  range  a  reasonable  accuracy  given  the difficulty of measuring ultrahigh temperatures. Both the  pyrometer  temperature and the current were recorded by a  computer. The system set up is described in greater detail  in Ref.  [11].  Two-stage  oxidation  experiments  were  performed  using  the  resistive  heating  system for  tracer  diffusion  studies.  Specimens  were  oxidized  for  a  predetermined  time under  flowing standard laboratory grade  16O2  (GTS Wilco, Allentown,  PA).  Then  the  specimen was  cooled  to  room temperature, which  happens  rapidly  once  the  current  flow ceases,  and  the  gas  flow was  turned  off.  The  specimen was  then  reinserted with  new copper wire  and  platinum foil,  as  occasionally  needed. The system 1.3 9 102 Pa  was  pumped  down  to  a  pressure  of  or  below,  then  current  and  temperature  acquisition  were  restarted,  and the  specimen was brought back up to temsystem was backfilled with 18O2 Sigma Aldrich, Saint Louis, MO)  perature. Then the  (97%  enrichment, 1 9 105 Pa  to  and the  specimen was oxidized for  a  second  predetermined time.  18O2  exposure  times were very short  (minutes or  less),  thus  the  stagnant  flow was  assumed to  have minimal  impact  on  the  testing.  Longer times “fogging” exposure was performed  are  not  possible with  stagnant  air  due  to  of  the  pyrometer window. A baseline using 16O2  for both oxidation stages  to show that  similar  oxidation  behavior was  observed whether  the  specimen  was heated once or  twice  (Figure 2). Only one  test  tem perature was performed due  to limited sample  supply.  In  addition,  Ref.  [11]  showed  a  test  comparing  oxidation  rate  and morphology  between  standard  box  furnace  and  resistive heating tests,  showing that  the  impact of  electric  current  through  the  diboride  has minimal  effect.  In addi tion,  the  ZrB2 to nickel)  has  a  conductivity  equivalent  to metals  (similar  and  it  is  therefore  unlikely  that  the  oxide  is  carrying  any  current. The  18O2  tracer  exposures  performed are  listed in Table 2. The  times  for  each stage  of  testing were  chosen in an attempt  to show both long and  short-term oxygen  diffusion  pathways, based temperatures.11,43  on  the  known  oxidation  rate  at  these  These  tests  showed  the  total  ZrO2 predicted  layer  growth  similar  to  or  smaller  than  the  growth,  demonstrating  that  cracks  resulting  in  short  circuit  oxidation  paths were  not  formed,  further  proving  normal  oxidation  behavior  for  resistive heating and double oxidation.  Conventional  mounting  and  polishing  of  specimen  cross-sections  resulted  in  surfaces  too  rough  for  time-of flight secondary ion mass spectrometry (ToF-SIMS), due to  the significant differences  in hardness of ZrB2, ZrO2, SiC, Instead, the double oxidization  C,  and  borosilicate  glass.  specimens were  fractured and half of  each specimen was  ion  polished  (Teledyne  Scientific,  Thousand Oaks, CA;  JEOL Cross  Section  Polisher  SM-09010,  Tokyo,  Japan;  University of Virginia, Hitachi  IM4000, Pleasanton, CA).  (A)  (B)  F I G U R E 2  Fracture cross-sections of ZrB2-30 vol% SiC oxidized using the resistive heating system at 1500°C (A)  for 20 minutes  consecutively and (B)  for 20 minutes in 2 stages (15 minutes and 5 minutes). Both cross-sections display the same morphology and similar  oxide thicknesses,  indicating no artifacts from double oxidation  T A B L E 2  Conditions for double oxidation experiments of ZrB2-30 vol% SiC  16O2  18O2  16O2 + 18O2  Temperature (°C)  Time (min)  Predicted ZrO2 growth (lm)†  Time (min)  Predicted ZrO2 growth (lm)a  Total measured ZrO2 growth (lm)  1500  19  11  1  <1  9.2 \\x06 1 8.6 \\x06 0.6 46.4 \\x06 6 80.4 \\x06 13  1500  10  8  9  3  1650  29  64  0.75  <1  1650  45  80  5  4  aZrO2 growth is calculated using experimental kp values.43  1768  |  CISSEL AND OPILA  \\x0c', 'CISSEL AND OPILA  |  1769  F I G U R E 3  outer surface of  ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1500°C for 19 minutes in 16O2 and 1 minute in 18O2. 18O is found on the indicates 20 lm  the borosilicate. Scale bar  For  this process, a thin layer of metal  (Au) was deposited  on the top surface to make contact with the knife blade that  defines the position of  the ion polished surface. The speci mens were  then  polished  back  directly  from the  fracture  surface with an Ar  ion beam. Due  to the  rough nature of  the  fracture  surface,  several  iterations of  ion milling were  performed. The  final  ion  polished  specimen  provided  the  smooth surface needed for ToF-SIMS.  Oxidized specimens were stored in a desiccator cabinet  to minimize  speci reaction of any B2O3 men with water vapor in the air. The oxidized specimens  remaining in the  (nonion  polished  sides) were  characterized  in  plan-view  and  cross-section  using  scanning  electron  microscopy  (SEM, FEI Quanta 600F, Hillsboro, OR)  and energy-dis persive MaxN  spectroscopy (EDS, Oxford Instruments Aztec X 150, Concord, MA)  after  the  surface was  coated  with  a  thin  layer  of  carbon.  (During  the  course  of  this  research, available SEM/EDS capabilities varied.) An oper ating voltage of 5 kV was used to maximize the EDS sig nal  intensity for  light elements.  F I G U R E 4  Line scan of 18O through oxide formed on ZrB2-30 vol% SiC exposed at 1500°C for 19 minutes in 16O2 and 1 minute in 18O2. Dashed lines mark approximate location of interfaces. 18O is  found in high concentrations at  the surface of  the borosilicate glass  \\x0c', 'Time-of-flight  secondary ion mass  spectrometry (ToF SIMS, PHI THRIFT V nanoToF, Chanhassan, MN;  ION TOF TOF SIMS V,  ION TOF,  Inc., Chestnut Ridge, NY)  was  performed  at  Case Western  Reserve  University  (CWRU)  and North Carolina State University (NCSU) on  cross-sectioned specimens after double oxidation in and 18O2. ToF-SIMS can distinguish between isotopes, and can therefore map the location of 16O and 18O. The ToF 16O2  SIMS  accelerates  an  ion  beam and  rasters  it  across  the  sample. These  primary  ions  sputter  the  surface,  ejecting  secondary ions  (or neutral  atoms), which are  separated by  mass. The secondary ions are accelerated into the detector  and analyzed by time of  flight, which varies  according to  mass-to-charge ratio.  There  are  two  data  collection modes  available, mass  resolution  and  spatial  resolution.  The  instrument  also  operates  in  either  positive  or  negative mode,  collecting  positive  or  negative  ions  favorably  in  each mode. Mass  resolution mode  has  a  resolution  on  the  order  of mil liatomic mass  units,  and was  used  here  to  confirm the  isotopes  present  in  the  oxidized  cross-section  specimen.  Mass  calibration was  conducted with  peaks  of CH, OH,  and C2H in negative operation mode and CH3, C2H3, and C3H5 in positive operation mode. The spatial resolution mode has a technical ~3 nm 0.5 lm without  resolution  between  and  fine  tuning,  and was  used  to map  the  location  of  the  ions  of  interest.  For  the  specimens  ana lyzed in this work, positive mode was used to collect Zr,  Si, and B, while 16O, and 18O.  negative mode was  used  to  collect C,  Limitations  of  the  ToF-SIMS  technique  include  the  requirement  for  a  flat  specimen to obtain data  as  surface  roughness  leads  to regions with low ion signal. Avoiding  epoxy in specimen mounting and polishing is also benefi cial  as  epoxy can lead to specimen charging. One of  the  systems used had an image  shift when switching between  positive  and  negative modes,  and  no  fine  stage  control.  This means  the  same  area  could not be  imaged for both  F I G U R E 5  SEM/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1500°C for 10 minutes in 16O2 and 9 minutes in 18O2. 18O is found throughout both the borosilicate glass and the ZrO2+C layer  F I G U R E 6  Line scan of 18O through the oxide formed on ZrB230 vol% SiC exposed at 1500°C for 10 minutes in 16O2 and 9 minutes in 18O2. Dashed lines mark approximate location of interfaces. 18O concentration is found to be relatively constant through the borosilicate glass. The 18O gradient in the ZrO2+C layer  may be an artifact of  the line scan sampling an irregular  interface  1770  |  CISSEL AND OPILA  \\x0c', 'modes, especially at high magnification (micrograph width ~100 lm). Due to image the exact  less  than  to  the  image  shift,  it was  not  possible  same grains  at high magnifi cation when  looking  for O (collected  in  negative mode)  and Si or Zr  (collected in positive mode), making it diffi18O was  cult  to  judge  if  the  in  the  borosilicate  or  the  ZrO2, unless EDS could be performed on the same grains. It should also be noted that different phases may release  ions at different  rates. This means  that  in one map,  inten sity changes between phases may not directly correspond  to  concentration  changes.  In  addition,  due  to  the  strong  gettering  nature  of ZrB2, both 16O, making  ZrB2 ZrO2 impossible to use the oxide. Since 18O2 (0.204% natural abun and  grains  appear  to  contain  it  16O  maps  to determine the location of  is  not  abundant  in  the atmosphere dance), maps of 18O are accurate location of oxide containing 18O.  representations  of  the  3  |  R E S U L T S  3.1  |  1500°C  ToF-SIMS results  for  a  specimen oxidized at 1500°C for  19 minutes in and 1 minute 18O was present at  16O2  in  18O2  are  given  in  Figure 3.  the surface of  the borosilicate  layer  and  not  observed  anywhere  else. This  figure  also  clearly shows the Si-depleted C-rich layer which forms at 1500°C.11 Figure 4 shows 18O2 with a only trace  a  line  scan of  the  sharp  peak  at  the  borosilicate  surface with  amounts lower  in the sample.  SEM/ToF-SIMS results for a specimen oxidized at 1500°C for 10 minutes in 16O2 and 9 minutes in 18O2 are given in Figure 5. 18O was present throughout both the borosilicate and ZrO2+C layer. Figure 6 shows a line scan 18O demonstrating relatively constant of the concentration  F I G U R E 7  SEM/EDS/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1500°C for 10 minutes in 16O2 and 9 minutes in 18O2, showing 18O throughout the ZrO2 grains even at the ZrO2/ZrB2 interface. Scale bar indicates 10 lm  CISSEL AND OPILA  |  1771  \\x0c', '1772  |  CISSEL AND OPILA  of 18O throughout the scale. Figure 7 compares ToF-SIMS 18O and EDS compositional mapping at 18O is  the same location  demonstrate  grains  found  in ZrO2  that  the  to  at  ZrO2/ZrB2 interface.  3.2  |  1650°C  SEM/ToF-SIMS results for a specimen oxidized at 1650°C for 29 minutes in 16O2 and 45 seconds in 18O2 are given in Figure 8. These maps clearly show 18O present at the gas/borosili cate glass  depletion layer  interface, as well as at the borosilicate glass/SiC interface. The 18O in the surface borosilicate  layer was not evenly distributed,  though SEM/EDS analysis  did not detect any microstructural/compositional variations between 18O rich areas and 16O areas. Figure 9 shows the region oxidized with 18O near the depletion layer, demonstrating the 18O was located in the borosilicate glass, which at  that  location was surrounded by ZrB2 grains. There was no evidence in this specimen of 18O in ZrO2 grains. This microstructure will be discussed in more detail later.  SEM/EDS/ToF-SIMS results for at 1650°C for 45 minutes in 16O2 in Figure 10 and Figure 11. As 1500°C (10 + 9 minutes)  second  given  the  are  a  specimen,  oxidized in 18O2 observed in 18O was  specimen  and 5 minutes  seen  throughout  the  borosilicate  layer  and  the  ZrO2  grains.  4  |  D I S C U S S I O N  4.1 | Oxygen diffusion pathways: Double oxidation experiments  4.1.1  |  1500°C  After oxidation at 1 minute in 18O2, surface of the borosilicate scale. This indicates that a small  ion is seen only on the exterior  19 minutes  the tracer  1500°C for  16O2  and  in  amount of network exchange has  observable permeation or  interface has occurred at  taken place but that no the ZrB2/SiC/ZrO2+C this time scale.  reaction at  F I G U R E 8  SEM/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1650°C for 29 minutes in 16O2 and 45 seconds in 18O2. 18O2 has exchanged with 16O2 in initial borosilicate scale, and permeated through the borosilicate glass and formed new Si18O2 at indicates 100 lm  SiC depletion layer  interface. Scale bar  the borosilicate glass/  \\x0c', 'After oxidation at 1500°C for 10 minutes  in 16O2 throughout  and  9 minutes  in  18O2,  the  tracer  ion  is  seen  the  borosilicate  scale.  This  result  indicates  that  either  net work diffusion in the borosilicate glass  is  responsible  for  oxygen transport or that process occurs. 18O is the ZrO2+C layer ZrO2 has also exchange  a  combined permeation/exchange  also  seen  in  the ZrO2 grains indicating that 16O in 18O.  in  (Figure 7),  the  exchanged with the 1500°C in ZrO2 was permeates through the  The  18O  rate  at  estimated,  assum ing  that  18O2  borosilicate  and  between  the  ZrO2 grains 1500°C,  quickly. After  19 minutes  of  oxidation  at  ZrO2 approximately  grains near ZrO2/ZrB2 3 lm in diameter. Assum18O needed to diffuse  the  interface were  ing the ZrO2 was fully exchanged, at least 1.5 lm. The  18O2  exposure was  9 minutes. The  diffusion  coefficient  in ZrO2 lower limit of the oxygen diffusion \\x0010 cm2/sec. This agreement with values extrapolated from nanocrystalline monoclinic ZrO2.30 AlterZrO2 grains at the ZrO2/ZrB2 interface from 18O2 been newly reacting with distinguish between new ZrO2 and 18O2 and old ZrO2 exchanged with a gradient in the 18O throughout  estimated  as  length  squared  over  time provides  a  rate  in ZrO2 in reasonable  on  the  order  of  10  value  is  Brossmann  for  natively,  these  could  have  formed  ZrB2. We by reacting ZrB2 18O. The lack of  cannot  formed  the  oxide  supports  rapid  exchange within  both  the  borosili cate  and ZrO2 grains.  4.1.2  |  1650°C  After oxidation at 45 seconds in 18O2, of the borosilicate layer and along the base material/oxide  1650°C for 29 minutes in and the 18O was found on the very outside  16O2  interface.  This  indicated  that  a  combination  of  network  exchange  at  the  surface  and  permeation  through  the  borosilicate glass occurs, for SiC oxidation.19,24,25 lichting’s oxygen diffusion 30%B2O3 \\x006 cm2/sec), (5 9 10 does allow for to have permeated through 100 lm of oxide in 45 seconds.44 The pres18O in  in  agreement with  the  literature  In addition, extrapolation of Sch1650°C  rate  for  to  18O2  ence  of  the  the  borosilicate  glass  at  the  interface  with the depletion layer  is consistent with this borosilicate  forming when the SiC grains at  the bottom of  the depletion  layer  actively  oxidized  to  SiO(g).  Figure 12  provides  a  schematic of 18O into the borosilicate formed just layer. Either 18O2 reacts with SiO(g) to form SiO2 via: Si16OðgÞ þ 1=218O2 ðgÞ ¼ Si16O18O  two  possible  reactions  for  incorporation  of  above  the  depletion  (6)  consistent with Reaction 4. Or  18O2  is  responsible  for  the  active oxidation reaction via:  18O2 ðgÞ þ SiC ¼ Si18OðgÞ þ C18O(g)  (7)  consistent with Reaction 2b. And the SiO(g)  subsequently  reacts with O2(g)  to form SiO2 (Reaction 4.).  F I G U R E 9  High-magnification SEM/EDS/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1650°C for 29 minutes in 16O2 and 45 seconds in 18O2, showing 18O in SiO2 at the no depletion/SiC depletion interface. Scale bar indicates 10 lm  CISSEL AND OPILA  |  1773  \\x0c', 'The  18O  distribution  in  the  borosilicate  glass  was  found  to  be  nonuniform after  both  the  short  and  long  18O2  exposures  (Figures 8  and  11). While  this  could  be  an  artifact  of  the  technique,  the  uniformity  of  the  total  ion  count map suggests that surface roughness does not the differing 18O intensities observed in some the 18O concentration in the borosilicate glass  account  for  locations.  If  is  actually nonuniform,  regions of  the  glass must  a have  higher  exchange  rate  of  oxygen. Karlsdottir  et al  suggest  compositional  variations  in  the  borosilicate  glass  occur  due  to  localized  formation  of B2O3  and  varying  ZrO2  solubility in borosilicate glass with B2O3 content.45 Variation in B2O3 content of the borosilicate glass has been oxide.46 shown already with  changes  of  depth  into  the  B2O3 on oxygen  content has also been shown to have a high impact glass.44  diffusion  rates  through  borosilicate  Thus, glass  regions higher  in B content would have more  ready  access  to  18O2  and would  likely  exchange with 18O-rich  18O2  earlier  than  silica-rich  regions. However,  regions of  the borosilicate glass were not to 18O-poor  found to contain  more B relative  regions within the  sensitivity  of EDS.  F I G U R E 1 0  SEM/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1650°C for 45 minutes in 16O2 and 5 minutes in 18O2. 18O was indicates 100 lm  found throughout  the borosilicate glass and ZrO2 grains. Scale bar  1774  |  CISSEL AND OPILA  \\x0c', 'By careful observation of  the intensities of 18O throughspecimen oxidized at 1650°C  out  the oxide  layers of  the  for  45 minutes  in  16O2  and  5 minutes  in  18O2  shown  in  Figure 13, a qualitative change in relative tracer concentra tion between the borosilicate glass and the ZrO2 phases noted. The ToF-SIMS map taken near the top of the oxide 18O intensity in the ZrO2 grains relaborosilicate glass. The ToF-SIMS map taken  is  scale shows a higher  tive  to  the  near the depletion region/bottom of the oxide scale shows a lower 18O intensity in the ZrO2 relative to the borosilicate glass. This indicates that the ZrO2 grains lower in the scale complete exchange than the ZrO2 18O2 exchange with the borosilicate glass. No  have  undergone  less  grains at  the surface. This  reveals  that  ZrO2 observation was made  is not  as  rapid as  that of  in  this  sample  of  the ZrO2/ZrB2 interface at high magnification to determine if new ZrO2 was grown by the oxidation of ZrB2 with 18O2. The O exchange rate at 1650°C in ZrO2 was also esti18O diffusion mated, again assuming to the edge of the rapid. After 50 minutes of oxidation at 1650°C,  grains  is  ZrO2 grains near the ZrO2/ZrB2 interface were approximately 5 lm in diameter, so 18O needed to diffuse at least 2.5 lm for the grain to be fully exchanged. The 18O2 exposure was 5 minutes. The diffusion coefficient in ZrO2 estitime provides a lower limit \\x0011 cm2/sec.  mated as  length squared over  to the diffusion rate on the order of 10  4.2  | Oxygen diffusion pathways  Three 18O in  possible mechanisms  can  explain the presence 18O could  of  the ZrO2. reacting ZrB2 counter  In  the  first  case,  diffuse  to  the  interface  through  the  ZrO2, vacancies  forming  new ZrO2 electrons.  by  diffusion  of oxygen 18O could  and  In  the  second,  the  diffuse  to  the  ZrB2 the glass  interface  through  fast  diffusion paths, for example, and form new 18O-enriched of 18O in the ZrO2 can exchange of 18O with 16O in ZrO2 that 16O2 oxidation step. Each of examined in light of the  phase  or  porosity  ZrO2. explained by the  Finally,  the  presence  be  was grown in the  first  stage  these  possibilities  is  F I G U R E 1 1  SEM/EDS/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1650°C for 45 minutes in 16O2 and 5 minutes in 18O2. The circle marks an area in 18O map showing less 18O in borosilicate glass. Scale bar indicates 20 lm  CISSEL AND OPILA  |  1775  \\x0c', 'observations  from experiments  conducted first  at 1650°C,  followed by a 1500°C results. At 1650°C, 18O for  discussion  of  the  observations  from the  the  thermally  grown ZrO2 the 5 minutes exposure, but not for  is  enriched  in  the 45 seconds  exposure.  It  is generally agreed that ZrO2 is an oxygen ion and growth of new dense oxide on a Zr-base  conductor  material is limited by the slower electronic species transport.12 This observation is based primarily on experiments  obtained  for Y2O3or CaO-stabilized ZrO2 ThO2-based systems.27,47,48 Two known exploring the conductivity of nominally pure oxides. Lasker and Rapp48  or  analogous  studies  exist  report  that nominally pure thoria is domi nated by ionic conductivity to oxygen partial pressures as \\x0028 atm or low as 10 temperatures as high as 1100°C. Vest and Tallan49 reported significant electronic contributions to nominally pure zirconia at 1400°C and oxygen partial pres\\x0014 atm, reversible  sures  below  10  however,  these  results  were  obtained without  electrodes  (unable  to  conduct  both electronic  and ionic  species). The  electronic  conduc tivity in ZrO2 increases relative to ionic conductivity as the dopant level decreases, the oxygen partial pressure decreases and as the temperature increases.47 Thus, a mixed  ionic/electronic conductor needed for oxidation of ZrB2 ZrO2 via ionic/electronic transport in ZrO2 is favored for high purity conditions (minimal impurity incorporation in low PO2’s and high temperatures. The purity  to  the ZrO2) at  of  the thermally grown ZrO2 in this experiment and the oxygen partial pressures across the \\x0020 atm at 1650°C from the gas the ZrB2/ZrO2 interface, respectively (as determined using the equilibrium module and FactPS database in FactSage).50 Thus, with the absence of conductivity data in the  is unknown  oxide  scale  span 1 to 10  interface  to  literature under  these  conditions,  it  is  surmised that  ionic  conductivity in the  thermally grown ZrO2 here dominates transport of oxygen vacancies through  and  oxidation  via  the ZrO2 is limited by low electronic conductivity. In the second mechanism, oxidation of ZrB2 18O-containing ZrO2 can occur by transport in fast paths such as pores or permeation through borosilicate glass. As the 18O has permeated the borosil to  form  seen in Figures 8 and 9,  icate glass phase to the oxidation interface even after only  45 seconds. New ZrO2 is then created at the ZrB2/ZrO2 interface incorporating 18O that has arrived via transport in  the borosilicate glass phase.  Finally,  in  a  third mechanism,  18O can be incorpo16O exchange  rated  in  the thermally 18O in  grown  ZrO2 glassy  via  with  the  the  adjacent  or  gas  phase.  It  is  concluded  from  Figure 10  that  18O  exchange has Zr18O2. reflect the  occurred with  the  preexisting  ZrO2 Thus, the estimated diffusion coefficients exchange—or self-diffusion—of  resulting  in  here  oxygen  in  this  thermally  grown  oxide.  It  is  seen  in  Figure 13  that  the  exchange  has  occurred more  readily  near  the  gas  interface  rather  than at the ZrB2/ZrO2 18O signal from ToF-SIMS. At 1500°C,  interface  from the  intensity of  the  a different oxide microstructure  is observed.  Thermal  oxidation  of ZrB2/SiC results layer and a sublayer of ZrO2 and carbon. After 1 minute of 18O2 exposure, exchange is only observed with the surface of the glassy network. After 9 minutes exposure in all enriched in 18O. Assuming again that growth of  in  a  glassy  surface  18O2,  oxides  are  new ZrO2 is not ZrO2 due to low electronic conductivity, the ZrO2 must be transporting through pores in the ZrO2 or carbon phases or in interfaces between the ZrO2 and the carbon phase. Pores were not observed in either phase at least  likely to occur by oxygen transport in dense the 18O observed in  at  the SEM length scale as shown in Figures 5 and 7. Thus,  the mechanism for inward oxygen transport 1500°C remains  to the ZrO2 scale be determined. Again, however, the  at  to  diffusion  coefficients  estimated  in  the  prior  section  likely  reflect  the exchange or  self-diffusion of oxygen in the ther mally grown ZrO2. Improving oxidation  resistance  of ZrB2-SiC is considered in future reusable  neces sary  for  this material  to be  applications.  This work  demonstrates  that  oxygen  trans port  in  both  the  borosilicate  and  the ZrO2 be utilized  is  rapid,  sug gesting multiple  strategies must  to  improve  oxidation resistance. Addition of  small amounts of  second  or  third  phases  have  been  shown  to  promote  the  F I G U R E 1 2  Schematic of 18O2 permeation through the borosilicate glass to form new Si18O2 at  the ZrO2/ZrB2 interface.  (In  the schematic, only the SiO2 component of  the borosilicate is labeled  since it  is the part  involved in the reactions.)  1776  |  CISSEL AND OPILA  \\x0c', 'formation of  immiscible glasses51-53 which can slow oxy gen  transport  in  the  glass  phase,  though  this  has only 1500°C.  been  demonstrated  at  temperatures  less  than  Oxygen  diffusion  in  ZrO2 strategies. Doping with  could  potentially 5+  be  slowed  utilizing  doping  cations  (Ta  and Nb) would  decrease  oxygen vacancies permeation.54  in  the ZrO2, However, this  thus  decreasing  oxygen  would potentially create oxides with lower melting points (Ta2O5 Tm = 1872°C, Nb2O5 Tm = 1512°C),55,56 decreasing the use temperature of the material. Thus, strategies  to improve  the oxidation resistance of ZrB2-SiC remain a challenge. While ZrB2-30 vol% SiC is limited by its for long-term applications  rapid  oxidation  at  ultrahigh  temperatures,  it  remains  of  interest  for  short-term single use applications.  A more rigorous  study with judicious choices of oxida tion  times  and  temperatures,  based  on  this  preliminary  work, would provide a more complete picture of  the oxy gen diffusion mechanisms.  In addition, use of  reduced oxy gen partial pressures will slow the oxidation rate and allow  for  testing  at  higher  temperatures  and  longer  times. This  low PO2 will which would be encountered in service.  also more  accurately  reflect  the  conditions  5  |  C O N C L U S I O N  ToF-SIMS  analyses  after  16O2-18O2  double  oxidation  experiments of ZrB2-30 vol% SiC have shown that oxygen exchange in the borosilicate glass begins rapidly at 1500°C, while  permeation  is  not  immediately  seen.  18O  exchange  and  diffusion  are seen in both the at 1500°C after  borosilicate  glass and the ZrO2 grains 1650°C, short exposure  9 minutes. At  to  18O2  after  16O2  oxidation  demonstrated that permeation and exchange in the borosili cate glass both occur. Oxygen exchange with the borosili cate glass may  be  compositionally dependent, but  further  work  is  required  to  confirm this. Oxygen  exchange was 1500°C  shown to be 1650°C. ZrO2 through fast diffusion  significant  in  ZrO2 likely  at  both  and  formation  is  facilitated  by  transport  paths  such as 1650°C.  permeation  in  the  borosilicate  glass  observed  at  The  fast  oxygen  F I G U R E 1 3  SEM/ToF-SIMS results for ZrB2-30 vol% SiC oxidized at 1650°C for 45 minutes in 16O2 and 5 minutes in 18O2, showing change of 18O contrast  from  brighter  in ZrO2 near surface to brighter  in  borosilicate near ZrO2/ZrB2 interface. The  total  ion count  is included to show overall  microstructural  influence  CISSEL AND OPILA  |  1777  \\x0c', 'transport  path  in  the  ZrO2/C  layer  formed  at  1500°C  remains to be identified.  A C K N OW L E D GM E N T S  Initial funding for this study was provided by the National Hypersonic Science Center - Materials and Structures. Dr.  Marshall  at Teledyne Scientific  and Dr. Robert Golden at  University of Virginia  assisted in the  ion polishing of  the  samples. The  authors would like  to thank Annette Marso lais (CWRU) and Chuanzhen Zhou (NCSU)  for  their assis tance with  the  ToF-SIMS,  as well  as  advice  on  data  interpretation. The authors would also like to acknowledge  Eric Neuman  and Dr. William Fahrenholtz  at Missouri  University  of  Science  and Technology  for  providing  the  ZrB2-30 vol% SiC material. This work was performed in part at the Analytical Instrumentation Facility (AIF) at  North Carolina State University, which is supported by the  State of North Carolina and the National Science Founda tion (award number ECCS-1542015). The AIF is a member  of  the North Carolina Research Triangle Nanotechnology  Network  (RTNN),  a  site  in the National Nanotechnology  Coordinated  Infrastructure  (NNCI).  Finally,  the  authors  would  like  to  thank Dr.  Triplicane  Parthasarathy, UES,  Inc.,  for helpful discussions.  O R C I D  Kathleen S. Cissel  http://orcid.org/0000-0003-0206-5476  R E F E R E N C E S  1. 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Oxygen diffusion mechanisms during high temperature oxidation of ZrB2-SiC. J Am Ceram Soc. 2018;101:1765-1779. https://doi.org/10.1111/jace.15298  CISSEL AND OPILA  |  1779  \\x0c']"
},{
  "_id": 206,
  "PDF": "Phase stability, hardness and oxidation behaviour of spark plasma sintered ZrB2-SiC-Si3N4 composites.pdf",
  "Text": "['Ceramics International 45 (2019) 9061-9073  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Phase stability, hardness and oxidation behaviour of spark plasma sintered ZrB2-SiC-Si3N4 composites  T  Sravan Kumar Thimmappaa, Brahma Raju Gollaa,∗, VV Bhanu Prasadb, Bhaskar Majumdarb, Bikramjit Basuc  a Metallurgical and Materials Engineering Department, National Institute of Technology, Warangal, 506 004, b Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad, India c Materials Research Center, Indian Institute of Science, Bangalore, 560012, India  India  A R T I C L E  I N F O  A B S T R A C T  Keywords: ZrB2 Spark plasma sintering Microstructure Densification Hardness Oxidation  Despite significant efforts to develop ultra-high temperature ceramics, the phase stability together with high hardness and oxidation resistance remains to be addressed in ZrB2-SiC based ceramics. ZrB2-20 vol% SiC (ZS20) ceramics with varying amounts of Si3N4 (2.5, 5 and 10 vol%) were processed by multi stage Spark Plasma Sintering (SPS) over a range of temperature (1800-1900 °C) for 3 min under 50 MPa. All the ZS20-Si3N4 composites could be densified to more than 98% theoretical density (ρth) after SPS at 1900 °C. The XRD, SEM-EDS analysis of the ZS20-Si3N4 composites revealed the presence of reaction product phases (ZrO2, BN, ZrN) along with SiC and ZrB2 major phases. Sintering reactions were proposed to explain the existence of such new phases and extinction of Si3N4. Thermo-Calc software was also used to further confirm the formation of these new phases in the ZS20-Si3N4 samples. The hardness of ZS20-Si3N4 composites varied between 25.50 and 30.56 GPa, in particular, ZrB2-20 vol% SiC-5vol% Si3N4 measured with the maximum hardness. In fact, it is the highest ever reported hardness for the ZrB2 composites. Considering oxidation resistance, the weight gain of ZrB2-20 vol%SiC increased (64-128 μm) after composites decreased (from 13.84 to 9.84 mg/cm2) and oxide layer thickness oxidation at 1500 °C for 10 h in air. The cross sectional microstructure of oxidized ZS20-Si3N4 composites consists of thick dense outer layer of SiO2, intermediate (ZrO2-SiO2) layer and unreacted bulk. The formation of dense SiO2 layer and absence of SiC depleted layer from the oxidized samples are signatures of improved oxidation resistance of Si3N4 reinforced ZrB2-20 vol% SiC.  1.  Introduction  Ceramics having melting point more than 3000 °C are popularly known as ultra-high temperature ceramics (UHTCs) and the UHTCs have drawn considerable research interest for the last one decade [1-4]. The UHTCs are potential candidate materials for thermal protection systems (TPS), components for hypersonic re-entry vehicles and for other structural and energy applications where corrosion-oxidationwear resistance are vitally important. Zirconium diboride (ZrB2) is one of the important class of boride UHTCs, in view of its unique combination of properties such as high melting point (3245 °C), high strength and elastic modulus, good electrical and thermal conductivity properties and high temperature oxidation resistance (up to 1200 °C) [2-4]. However, obtainment of fully dense ZrB2 is difficult due to its strong covalent bonding, low self-diffusion coefficient and existence of oxides (ZrO2 and B2O3) on the starting powders, which essentially requires  very high sintering temperatures and long holding times [4]. The use of sintering additives and advanced sintering techniques have been attempted to improve sinterability, to refine the microstructure and to enhance the mechanical and oxidation properties [5,6]. Both conventional (pressureless sintering, hot press) and advanced sintering (Spark plasma sintering (SPS), Flash sintering, Microwave sintering) techniques have been employed to densify ZrB2 materials. Although metallic and non-metallic sintering additives were used for enhancing densification of ZrB2, the non-metallic sintering additives are [4,7-26]. advantageous in improving high temperature properties In particular, use of SiC (5-50 vol%) as the additive for ZrB2 has been extensively studied as it is effective in enhancing both room and elevated temperature properties of ZrB2 [2]. In fact, addition of 20 vol% SiC to ZrB2 was proved to be optimal composition for hypersonic applications as the ZrB2 composite exhibited good thermo-mechanical properties [2,27,28]. Hence, in this work, ZrB2-20 vol% SiC (ZS20) was  ∗ Corresponding author. E-mail addresses: gbraju121@gmail.com, gbraju@nitw.ac.in (B.R. Golla).  https://doi.org/10.1016/j.ceramint.2019.01.243 Received 20 December 2018; Received in revised form 25 January 2019; Accepted 29 January 2019  Available online 30 January 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c', \"S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  selected as a base composition for developing the composites. The oxidation of monolithic ZrB2 in air is more prominent at high temperatures (more than 1200 °C) and it leads to the formation of ZrO2 and B2O3 phases on its surface. However, B2O3 evaporates in the temperature range of 1100-1400 °C and porous ZrO2 left alone, which lowers the oxidation resistance of ZrB2 [5,8]. The oxidation resistance of ZrB2 can be improved by adding different additives such as MoSi2 [7,18,20], TaSi2 [21-23], ZrC [24-26], AlN [8], ZrO2 [9,10] and SiC [1,11-13] etc. as they form protective oxide layer. In case of ZrB2-SiC, the formation of protective borosilicate glass or SiO2 rich oxide layer on the surface inhibits the diffusion of oxygen and improves the oxidation resistance [14,15]. As far as the novelty of the present work is concerned, multi stage spark plasma sintering (SPS) was employed to densify ZrB2-20 vol% SiC-Xvol.% Si3N4 (X = 2.5, 5 and 10) composites. Limited research is available on Si3N4 reinforced ZrB2 ceramics and also use of multi stage SPS in developing boride ceramics [16,27-29]. The combined effect of SiC and Si3N4 on ZrB2 has not been explored much. In fact, Si3N4 is one of the important silicon based ceramics that has potential for high temperature applications. In the light of this, an attempt has been made to find whether ZrB2 ceramics could be densified at low sintering temperatures? Efforts were made to understand underlying sintering mechanisms of ZrB2SiC-Si3N4. The microstructure, hardness and oxidation properties of ZrB2 were evaluated to realize the combined effect of SiC-Si3N4 addition. The multi stage SPS experiments were carried out over a range of temperature 1800-1900 °C for a short duration of 3 min under the uniaxial pressure of 50 MPa. Isothermal oxidation tests were performed at 1500 °C for a duration of 10 h to assess thermal stability of the ZrB2 composites. The microstructural studies were performed using XRD, SEM-EDS and Thermo-Calc software.  2. Experimental  In this work, commercially available ZrB2 (H.C. Starck Grade B, GmbH and Co., Goslar, Germany), SiC (purity > 99.8%, Alfa Aesar) and Si3N4 (purity > 99.8%, Alfa Aesar) powders were used for the processing of ZrB2-SiC-Si3N4 ceramic composites. To prepare ZrB2-20 vol% SiC-Xvol.% Si3N4 (X = 2.5, 5 and 10) composites, the powders in appropriate amount were mixed by wet ball-milling using a planetary ball mill (Fritsch Pulverisette 6, Germany). For convenience, the base composition ZrB2-20 vol% SiC (ZS20) reinforced with 2.5, 5.0 and 10.0 vol% Si3N4 was designated as ZS20-2.5SN, ZS20-5SN and ZS2010SN, respectively. The ball milling was carried out at 200 rpm for 6 h using Si3N4 milling media (the Si3N4 balls and the vials were lined with Si3N4) and toluene was used as a liquid medium with balls to powders weight ratio as 2:1. The use of toluene prevents oxidation and also acts as a process control reagent by maintaining lower temperature during milling. After mixing, the slurry was dried in a rotary vacuum evaporator at 98 °C to remove toluene and minimize the particles agglomeration. The dried powders were sintered at different stages (multi stage) using SPS (Model: SPS 25-10, GT Advanced Technologies, USA) at temperature of 1900 °C for 3 min under 50 MPa pressure in a vacuum (10−3 Pa) using graphite die and punches. The multi stage SPS schedule is presented in Fig. 1. The samples of 15 mm in diameter and thickness of 3-4 mm were prepared with the SPS. After the SPS, the specimens were subjected to belt grinding, polishing and ultrasonic cleaning (using acetone). The final density of samples was measured using Archimede's method, while the theoretical density of ZS20-Si3N4 composites were calculated by rule of mixture. The crystalline phases in the starting powders, ball milled compositions and SPSed samples were characterized by X-ray diffraction (XRD, PANalytical X'Pert Pro, Holland, CuKα = 1.5405 Aº). The microstructure of starting powders and sintered ZrB2 composites was carried out using a scanning electron microscope (SEM: TESCAN VEGA 3 LMU). The elemental analysis of ZrB2 samples was accomplished by energy dispersive spectroscopy (EDS: Oxford Instruments) attached to SEM. The ZrB2 samples were  Fig. 1. A schematic showing heating-cooling cycle followed while using multi stage SPS to densify ZrB2-20 vol% SiC-Xvol.% Si3N4 (X = 2.5, 5 and 10) at a pressure of 50 MPa under vacuum.  polished with a diamond paste up to a 1 μm finish. The microstructure of samples was revealed by chemical etching of the samples in 10 HCl-1 HNO3 solution at room temperature for 40 s. The average grain size of ZrB2 was measured on multiple SEM images with ImageJ software package (ImageJ 1.51j8, National Institute of Health, USA). Further the phases formed during sintering were confirmed using Thermo-Calc (2017b) software. The hardness of the samples was measured under 2 kg load at dwell time of 15 s using Vickers hardness tester (M/s. Shimadzu, HMV, Japan). At least 5 hardness measurements were taken for each individual sample. The oxidation studies were characterized by measuring weight changes of the ZrB2 composites (rectangular bars of 2 × 4 × 8 mm3 were cut from initial cylindrical samples), which were heated in stagnant air at 1500 °C for 10 h in a box type furnace (Nabertherm, MoSi2 heating element). Prior to oxidation, the specimens were thoroughly cleaned using acetone in an ultrasonic bath. The cleaned specimens were placed in a pure alumina crucible with minimal contact area and then the crucible was inserted at the center portion of furnace. Detailed microstructural characterization of oxidized samples (both surface and cross-section) were evaluated using XRD and SEMEDS.  3. Results and discussion  3.1. Microstructural characterization  The morphology of starting powders that are used in preparing ZrB2-SiC-Si3N4 ceramic composites is presented in Fig. 2. The ZrB2 particles are of platelet shape and its particle size observed to vary in the range of 0.5-4.0 μm. Whereas the SiC and Si3N4 powders are of much finer, irregular and in the agglomerated form (Fig. 2b and c). The particle size of SiC existed in the range between 0.3 and 2.6 μm and Si3N4 particles size varied from 0.1 to 1.6 μm. Fig. 3a shows the XRD patterns of starting powders. From the XRD, it can be noticed that the powders consisted of its characteristic peaks of individual materials (ZrB2, SiC and Si3N4) and no other additional phases could be identified within the detection limit of XRD. The XRD phases of ZrB2-SiC-Si3N4 powder compositions after ball-milling (at 200 rpm for 6 h) is also shown in Fig. 3b. In case of ZS20-2.5SN, only ZrB2 and SiC phases could be observed and no traces of Si3N4 could be detected. It may be due to the detection limit of XRD to identify small fraction (2.5 vol%) of Si3N4. On the other hand, the presence of ZrB2, SiC and Si3N4 phases were very evident in ZS20-5SN and ZS20-10SN samples. It also can be noted that the peak intensity of Si3N4 increased with increasing addition of Si3N4  9062  \\x0c\", \"S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 3. (a) XRD patterns of ZrB2, SiC and Si3N4 starting powders and (b) ballmilled ZrB2-20 vol% SiC-Si3N4 powder compositions. Ball milling was carried out at 200 rpm for 6 h by maintaining ball to powders weight ratio at 2:1.  SiC phases along with new phases (BN, ZrO2 and ZrN) without any hint of Si3N4. Presence of such new phases in the SPSed ZrB2 composites clearly indicate involvement of sintering reactions. A careful look at XRD patterns reveals that the peak intensity of BN and ZrO2 phases decreases, while the ZrN peak intensity increases with increasing the addition of Si3N4 to ZrB2-20 vol% SiC. It is interesting to note that the crystalline ZrO2 phase could not be identified in the starting or ball milled powders, however, the phase is noticed in sintered (SPS processed) ZrB2 composites. Similar kind of observations (i.e. absence of ZrO2 phase in powders and its presence in sintered samples) was also reported for hot pressed ZrB2-MoSi2 and ZrB2-NdB6 composites [17,19]. In the following, more details will be discussed.  Fig. 2. SEM of starting powders of (a) ZrB2, (b) SiC and (c) Si3N4.  3.2. Sintering reactions  amount to ZS20 powder compositions. The phase evolution of ZS20-XSN (X = 2.5, 5.0 and 10.0 vol%) composites after multi-stage spark plasma sintering (Temperature: 1900 °C, Pressure: 50 MPa, Time: 3 min) under vacuum is presented in Fig. 4. Interestingly, all the ZS20-xSN composites consists of ZrB2 and  From the XRD patterns, it is clear that the SPSed ZrB2 composites consist peaks of ZrN, BN and ZrO2 along with the ZrB2 and SiC phases (Fig. 4). No traces of Si3N4 could be observed in the XRD spectra of assintered ZrB2-SiC-Si3N4 ceramics. It might be due to complete consumption of Si3N4 while it's reaction with the other phases during the SPS process. It was reported that the oxide impurities (B2O3 and ZrO2)  9063  \\x0c\", 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 4. XRD phase analysis of ZrB2-SiC-Xvol.% Si3N4 (X = 2.5, 5 and 10) after multi stage SPS at 1900 °C, 50 MPa for 3 min.  present as a thin layer on the outer surface of ZrB2 powder particles [1]. At high temperatures it is possible that the Si3N4 ceramic can decopose into Si and N2 [29]. The sintering reactions are possible to take place due to the presence of impurities in the powders, decomposition of Si3N4 and presence of carbonaceous medium during SPS. In fact, the SPS of ZrB2 composites was carried using graphite dies. Hence, the ZrN and BN phases might have formed through the chemical reactions between the oxide impurities present on the surface of ZrB2 powders and Si3N4 or Si, N2 and C etc. according to following reactions.  2B2O3 + Si3N4 = 4BN + 3SiO2  3ZrO2 + Si3N4 = 3ZrN +3SiO2 +0.5N2  Si3N4 = 3Si + 2N2  2ZrB2 + 3Si + 3N2 + 3C = 2ZrN +4BN + 3SiC  3ZrB2 + 2Si + 4N2 = 2ZrN + 6BN + ZrSi2  ZrB2 + 1.5N2 = ZrN + 2BN  (1)  (2)  (3)  (4)  (5)  (6)  To understand the thermodynamic feasibility of the above reactions, the free energy (ΔG) formation of the reactions (1-6) as a function of temperature is presented in Fig. 5a. It can be observed that reactions 1, 3, 4 and 6 are thermodynamically feasible as the free energy of these reactions are negative and the other reactions (2 and 5) do not take place as they exhibit positive free energy of formation. From Fig. 5a, it can be observed that reaction (1) is strongly negative at the sintering temperature of 1900 °C. So, the BN formation may be possible by reaction (1). Monteverde et al. [16] also reported BN formation for ZrB22.5Si3N4, which was consolidated via hot press at 1700 °C. By closely observing Fig. 5a, it can be realized that reaction 1, 3 and 4 are more favorable during SPS process due to its strong negative free energy at 1900 °C. These sintering reactions indicate consumption of Si3N4 by forming new phases ZrN, BN, SiC and SiO2. However, SiO2 phase could not be observed from XRD or SEM analysis of the samples. Fig. 5b shows the isothermal section of Zr-N-B ternary phase diagram at 1900 °C, that is constructed using thermocalc software. The phase diagram shows that ZrB2, BN and ZrN phases are thermodynamically compatible. Based on this ternary phase diagram and proposed sintering reactions, it is very clear that formation of ZrN and BN phases are likely in ZrB2-SiC-Si3N4 composites at a sintering temperature of 1900 °C. As explained earlier, the XRD results also corroborate the presence of ZrN and BN phases. Talmy et al. reported the formation of BN, ZrSi2, ZrN phases along with ZrB2 in ZrB2-(5-35)vol%  9064  Fig. 5. (a) Gibbs free energy change of proposed sintering reactions (1-6) as a function of temperature and (b) isothermal section of Zr-N-B ternary phase diagram (at 1900 °C) constructed using Thermo-Calc software.  Si3N4 ceramics (hot press:1850 °C, 20 MPa, 1 h) [27]. Bellosi and Monteverde observed BN, ZrO2 and glassy phase (B-N-O-Zr-Si) for hot pressed ZrB2-5vol% Si3N4 composite [28]. They attributed the formation of such phases to reaction between Si3N4 and oxides of boron and silicon oxides present on ZrB2 and Si3N4 powders and ZrB2. In another work, Monteverde et al. [16] proved the presence of grain boundary phases (BN, ZrO2, ZrSi2 and B-N-O-Zr-Si glassy phase) with thermodynamic calculations for ZrB2-2.5 wt% Si3N4 composites. Similarly, Park et al. observed TiN, BN and amorphous SiO2 for hot pressed TiB2Si3N4 ceramics due the sintering reaction between the Si3N4 and TiO2 that was present on the surface of TiB2 powders [30]. Guo et al. reported the formation of ZrN, BN, ZrSi2, SiC or Si phases for Si3N4-30 wt % ZrB2 composite that were hot pressed at 1700 °C under Ar or N2 atmosphere [29]. In case of in-situ hot pressed ZrB2-20 vol% SiC-6.05 vol % ZrC composite, no additional secondary phases were reported [31]. Interestingly, presence of BN nano-platelets, TiN and crystalline SiO2 phases were reported for SPSed TiB2Si3N4 due to the reaction between Si3N4 and TiO2 and B2O3 oxides that were present on TiB2 powders [32].  \\x0c', '3.3. Densification  The heating schedule that is followed for SPS of the ZrB2 composites is presented as a schematic in Fig. 1. As it is well known, the SPS is a very fast sintering process in consolidating the materials, in particular, the whole sintering process lasts within a very short time (mostly below 60 min). The various factors such as high heating and cooling rates, joule heating, electric discharge in SPS is believed to aid in the densification and also obtainment of fine microstructure due to enhanced sintering kinetics [33]. From Fig. 1, it should also be apparent that a multi stage heating schedule was followed during the SPS process. Table 1 records the processing conditions together with density and hardness measurements of SPSed ZrB2 composites. It should be noted here that we have presented theoretical density (ρth) (based on initial powder compositions) of the ZrB2 composites using rule of mixture by considering theoretical density of each individual phase. To calculate theoretical density of ZrB2 composites, the densities of ZrB2, SiC and Si3N4 were taken as 6.08 g/cc [34], 3.21 g/cc [7] and 3.2 g/cc [35], respectively. In the literature, most commonly the researchers presented densification of ZrB2 materials based on the initial powder compositions despite the variation in phases of sintered samples final microstructure. We believe such measurements may not be true representation of densification, if new phases are involved in the sintered the relative density (%ρth) of ZS20-Si3N4 composites samples. Hence, were alternatively estimated by determining the porosity of the samples from the SEM micrographs of sintered samples using ImageJ software. The preliminary SPS experiments revealed that the measured experimental density of ZS20-2.5SN increases from 5.31 to 5.37 g/cc with increasing SPS temperature from 1800 to 1900 °C. Hence, the other ZrB2 composites were also SPSed at 1900 °C with a view to achieve high density. A common observation is that the density of ZrB2-20SiC composites decreased (5.37-5.11 g/cc) with increasing addition of Si3N4, since its density is considerably low compared to ZrB2. Almost full density (> 98%ρth) of ZrB2 composites was achieved when Si3N4 content was more than or equal to 2.5 vol%. The relative density was slightly lowered when higher amount (≥5 vol%) of Si3N4 was added to ZrB2. It may be due to the formation of more amount of secondary phases in the ZrB2 composites. However, the microstructure of SPSed ZrB2 samples evidences almost complete densification of ZrB2-SiC-Si3N4 composites as there were no indication of porosity (see Fig. 6). Bellosi and Monteverde studied the effect of different additives (Ni, SiC, Si3N4, AlN) on the densification, mechanical and oxidation properties of hot pressed ZrB2 [28]. The hot press (HP) experiments were performed over a range of temperature (1700-1870 °C), for 10-30 min under 30 MPa pressure. In case of monolithic ZrB2, a maximum density of 87%ρth was obtained even after HP at 1870 °C. On the other hand, the ZrB2-5vol% Si3N4 and ZrB2-19 vol% SiC composites could be densified to 98%ρth. Ahmadi et al. reported obtainment of 95%ρth for ZrB2-30 vol% SiC after HP at 1900 °C under 10 MPa for 120 min [36]. The ZrB2-30 vol% SiC composite could be fully densified with the addition of Si3N4 (up to 5 vol%). The back scattered electron (BSE)-SEM images indicate presence of different contrasting (bright, grey and dark) phases in SPSed ZrB2-SiCSi3N4 composites (Fig. 6). The SEM-EDS confirms bright contrasting phase as ZrB2, the grey phase as SiC and the dark phase as ZrN. The SiC and ZrN phases were well dispersed within the microstructure. The grain size of ZrB2 varied narrowly between 3.37 and 3.61 μm and it slightly increased with increasing addition of Si3N4. In fact, the grain size of sintered ZrB2 is very much comparable with initial ZrB2 powders particle size and clearly indicate insignificant grain gorwth for the SPSed ZrB2 composites. As discussed earlier, the sintering reactions represent involvement of liquid phase during sintering. Nevertheless, the microstructure of multi-stage SPS of ZrB2 results in unifirm microtsructure without any signs of abnormal grain growth, which is a typical characteristic of liquid phase sintered materials. From this it can be understood that the densfication of SPSed ZrB2-SiC-Si3N4 composites  T  a  b  l  e  1  D  e  n  s  i  fi  c  a  t  i  o  n  ,  m  i  c  r  o  s  t  r  u  c  t  u  r  e  ,  a  d n  h  a  r  n d  e  s s  o  f  S  P  S  e  d  Z  r  B  2   S  i  C   S  i  3  N  4  c  o  m  p  o  s  i  t  e  s  .  C  o  m  p  o  s  i  t  i  o  n  S  i  n  t  e  r  i  n  g  c  o  d n  i  t  i  o  n  s  (  °  C  ,  M  P  a  ,  m  i  n  )  ρ  t  h  g  /  c c  (  u  s  i  n  g  s  t  a  r  t  i  n  g  p  o  w  d  e  s r  )  E  x  p  .  D  e  n  s  i  t  y  (  g  /  c c  )  R  D  (  f  r  o  m  s  t  a  r  t  i  n  g  p  o  w  d  e  s r  )  R  D  (  f  r  o  m  S  E  M  i  m  a  g  e  s  )  G  r  a  i  n  s  i  z  e  (  μ  m  )  H  v  (  G  P  a  )  h P  a  s  e  s  Z  S  0 2   2  .  5  S  N  0 0 8 1  ,  0 5  ,  3  5  .  3 4  5  .  1 3  7 9  .  8  7 9  .  2 9  3  .  3 2  ±  0  .  7 6  0 2  .  1 6  ±  1  .  6 7  Z  r  B  2  ,  S  i  C  ,  B  N  ,  Z  r  N  ,  Z  r  O  2  Z  S  0 2   2  .  5  S  N  0 5 8 1  ,  0 5  ,  3  5  .  3 4  5  .  3 3  8 9  .  2  8 9  .  6 5  3  .  7 3  ±  0  .  2 6  5 2  .  9 5  ±  1  .  8 3  Z  r  B  2  ,  S  i  C  ,  B  N  ,  Z  r  N  ,  Z  r  O  2  Z  S  0 2   2  .  5  S  N  0 0 9 1  ,  0 5  ,  3  5  .  3 4  5  .  7 3  8 9  .  8  8 9  .  9 7  3  .  7 3  ±  0  .  0 1  5 2  .  0 5  ±  1  .  6 7  Z  r  B  2  ,  S  i  C  ,  B  N  ,  Z  r  N  ,  Z  r  O  2  Z  S  0 2   5  S  N  0 0 9 1  ,  0 5  ,  3  5  .  6 3  5  .  8 2  8 9  .  5  8 9  .  9 2  3  .  4 4  ±  0  .  1 1  0 3  .  6 5  ±  2  .  0 3  Z  r  B  2  ,  S  i  C  ,  B  N  ,  Z  r  N  ,  Z  r  O  2  Z  S  0 2   0 1  S  N  0 0 9 1  ,  0 5  ,  3  5  .  1 2  5  .  1 1  8 9  .  0  8 9  .  3 7  3  .  1 6  ±  0  .  2 1  5 2  .  9 4  ±  2  .  8 8  Z  r  B  2  ,  S  i  C  ,  B  N  ,  Z  r  N  ,  Z  r  O  2  S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  9065  \\x0c', \"S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 6. Microstructures of (a) ZS20-2.5SN, (b) ZS20-5SN and (c) ZS20-10SN samples after multi stage SPS at 1900 °C, 50 MPa for 3 min and the corresponding EDS of various phases observed in the samples i.e., (d) ZrB2 (bright phase), (e) SiC (grey phase) and (f) ZrN (dark phase).  occurred mainly by Si3N4 addition, sintering reactions and liquid phase sintering. The presence of oxide impurity reportedly hinders the densification of borides due to evaporation-condensation and grain gorwth mechanisms. Since Si3N4 addition lead to sintering reactions, it is consumed either due to it's decomposition at high temperature or via sintering reactions in eliminating the oxide impurity of ZrB2. The  addition of Si3N4 not only removes the oxide impurities from the surface of ZrB2 and it controls the grain growth of ZrB2 by forming nitride phases (ZrN, BN) which hinders the grain growth of ZrB2. Hence, Si3N4 is an effective sintering additive in enhancing densification with uniform fine microstructure of SPSed ZrB2 composites. The secondary electron images of ZS20-10SN composite is presented in Fig. 7. The  9066  \\x0c\", 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 7. (a) Low and (b) high magnification SE images of ZS20-10SN sample, showing growth of BN platelets on the surface.  Fig. 8. SEM micrographs of represented by circles.  fracture surfaces of (a) ZS20-2.5SN, (b) ZS20-5SN and (c and d) ZS20-10SN. The presence of nano sized needle shaped platelets were  9067  \\x0c', 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Table 2  Weight change and free energy formation (ΔG) of ZrB2-SiC-Si3N4 composites with different oxidation reactions.  S. No.  Reaction No.  Oxidation Reactions  Mass per mole of  1 2 3 4 5 6  7 8 9 10 11 12  ZrB2(s)+5/2O2(g) = ZrO2(s)+B2O3(l) B2O3(l) = B2O3(g)  SiC(s)+3/2O2(g) = SiO2(l)+CO(g)  SiC(s)+O2(g) = SiO(g)+CO(g) 2ZrN(s)+2O2(g) = 2ZrO2(s)+N2(g) 4BN(s)+3O2(g) = 2B2O3(l)+2N2(g)  ZrB2 B2O3 SiC SiC ZrN BN  Δw, g  80 −69.6 20 −40 36 40  ΔG (KCal/mol) at oxidation Temp. (1500 °C)  −320.07 22.11 −192.82 −109.19 −272.40 −318.69  graphene structure, which formed due to 2D hexagonal growth of BN [36]. The microstructure of fracture surfaces of SPSed ZrB2-based composites is shown in Fig. 8. Mixed mode (transgranular and intergranular) of fracture without any indication of pores was noticed for ZS20 with 2.5, 5 and 10 vol% Si3N4 composites. The grains in the fracture surface were completely connected together, which signifies good densification of samples. A close observation of Fig. 8 clearly indicates the presence of nano sized needle shaped platelets (represented by circles).  3.4. Hardness  The hardness of SPSed Si3N4 reinforced ZS20 composites is presented in Table 1. The hardness of ZS20-2.5SN improved significantly from 20.61 to 25.59 GPa with increasing sintering temperature from 1800 to 1850 °C. Such improvement in hardness can be attributed to enhanced densification of the sample. Nevertheless, the enhancement in densification is not as drastic as the hardness of the ZS20-2.5SN composite. On the otherhand, further increasing the sintering temperature (up to 1900 °C) did not yield any notable improvement in the hardness or density of the composite. Interestingly, further increasing the amount Si3N4 (up to 5 vol%) resulted outstanding enhancement of hardness (30.56 GPa) and densification (more than 98%ρth). However, the hardness was reduced to 25.49 GPa with ZS20-10SN due to its relatively low density and more amount of secondary grain boundary phases. Bellosi and Monteverde reported maximum hardness of 13.4 GPa for the hot pressed ZrB2-5vol% Si3N4 composites, which was  Fig. 9. Weight gain and oxide layer after oxidation at 1500 °C for 10 h.  thickness as a function of Si3N4 amount  presence of needle like BN whiskers can be noticed from the micrographs. It is possible that BN whiskers grow during sintering. Similarly, the presence of BN platelets was also observed for hot pressed ZrB230 vol% SiC-(1-5 vol%) Si3N4 and SPSed TiB2-5wt.% Si3N4 systems [32,36]. Ahmadi et al. reported that the BN platelets resemble like  Table 3  Comparison of processing conditions, microstructure and oxidation characteristics of different ZrB2based composites.  Composition (vol%)  Sintering conditions (°C, MPa, min)  RD  Microstructure of sintered ZrB2 samples  Oxidation conditions (Temp.°C, Time h)  Weight gain (mg/cm2)  Oxide layer thickness (μm)  Oxide phases  Ref.  ZS20-2.5SN ZS20-5SN ZS20-10SN ZrB2 ZrB2 ZrB2-5 Si3N4 ZrB2-5 Si3N4 ZrB2-25 SiC ZrB2-20 SiC15 Graphite ZrB2-20SiC-40MoSi2  1900, 50, 3 1900, 50, 3 1900, 50, 3 HP, 1850, 20, 60 HP, 1850, 20, 60 HP, 1850, 20, 60 HP, 1850, 20, 60 HP, 1850, 20, 60 HP, 1900, 30, 60  HP, 1800, 30, 30  ZS20  SPS, 2100, 7, 5  HP, 2000, 25, 60 HP, 2000, 30, 60 HP, 2000, 30, 60 RHP, 1800, 20, 60 followed by hot forging HP, 1950, 32, 45 SPS, 2000, 32, 30 HP, 1900, 30, 60  ZrB2-5wt.% B4C ZrB220SiC ZrB220SiC-5AlN ZrB2-20 MoSi2  ZrB2-30 SiC ZrB2-30 SiC ZrB2-20 SiC-20 ZrC  Laminated BN/ZrB2SiC ZrB2-10SiCf-10ZrSi2  99 > 98 > 98 > 99   99  99.8 > 98  > 98 > 95  ZrB2, SiC, BN, ZrO2 ZrB2, SiC, BN, ZrO2, ZrN ZrB2, SiC, BN, ZrO2, ZrN ZrB2 ZrB2 ZrB2, ZrSi2, BN, ZrN ZrB2, ZrSi2, BN, ZrN ZrB2, SiC ZrB2, SiC, C, ZrC  1500, 10 1500, 10 1500, 10 1300, 2 1500, 2 1300, 2 1500, 2 1500, 2 1800, 0.5, 2000 Pa (PO2)  13.84 13.28 9.84 6.0 25.0 14.0 28.0 5.0 22.0  ZrB2, SiC, MoSi2  1500, 10  ZrB2, SiC  ZrB2, B4C ZrB2, SiC ZrB2, MoSi2  ZrB2, SiC ZrB2, SiC ZrB2, SiC, ZrC  2200, 20 S in (30 vol% H2O+ 70 vol% Ar) 1500, 5 1800, 1 1800, 1 1500, 12  1500, 0.5 2000, 0.00277 1600, 4  64 94 128 160 140  291  775 155  13 30 65   65  SiO2, ZrO2 SiO2, ZrO2 SiO2, ZrO2 ZrO2, B2O3 ZrO2 ZrO2 ZrO2 SiO2, ZrO2 mZrO2, tZrO2 SiO2, ZrO2, MoB, ZrSiO4 SiO2, ZrO2  ZrO2 SiO2, ZrO2 SiO2, ZrO2, MoB ZrO2, SiO2 mZrO2 SiO2, ZrO2, ZrCxOy SiO2, ZrO2  SiO2, ZrO2  pw pw pw [27] “ “  \" \" [7]  [20]  [11]  [40] [41] “  [42,43]  [44] [45] [46]  [47]  [48]  3.0   16.2 15.0 57.0 7.0  0.5 55.0  10.0  12.7  HP, 1900, 30, 60   ZrB2, SiC, BN  1500, 10  HP, 1600, 50, 10  100  ZrB2, SiC, SiO2, ZrO2, ZrSi2  1500, 10  9068  \\x0c', 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 10. Cross sectional SEM micrographs of oxidized (a, d, e) ZS-2.5SN, (b, f, g) ZS-5SN and (c, h, i) ZS-10SN and its high magnification images of ZrO2-SiO2 layer (d, f and h) and unreacted bulk (e, g and i).  densified to 98%ρth [28]. A slight increase in hardness of 14.2 GPa was noticed for ZrB2-19 vol% SiC. Gupta et al. reported that ZrB2-SiC-TiSi2 could exhibit high hardness (27 GPa) for samples subjected to multi stage SPS when compared to single step conventional SPS due to its refined microstructure [34]. Hence, in multi-stage SPS the composites would achieve homogeneous microstructure and high density with fine grain size, which intended to enhance the mechanical properties. A close look at Table 2 reveals that there is no much difference in the grain size of the ZS20-SN composites. The grain size of ZrB2 slightly increased from 3.37 to 3.61 μm with inreasing addition of Si3N4. Since the densification and grain size of all the ZrB2 composites (SPS at 1900 °C for 3 min) are almost in the same range, the difference in hardness of the samples can be mainly attributed to the amount of the phases that are present in it and its distribution. It reflects that when higher amount of secondary phases present, the hardness of ZrB2 decreases. Park et al. [30] also observed similar kind of behaviour for TiB2-Si3N4 composites. Based on the present results, it can be realized  that ZS20-5SN is optimal composition among all as it exhibited maximum hardness and density. The hardness of ZS20 with 5 vol% of Si3N4 was improved due to enhancement in density and uniform distribution of secondary phases.  3.5. Oxidation of ZrB2-20SiC-(2.5-10) Si3N4  Fig. 9 presents the specific weight gain and oxide layer thickness of ZrB2-20 vol% SiC-Xvol.% Si3N4 (X = 2.5, 5.0 and 10.0) composites after oxidation at 1500 °C for 10 h. As the amount of Si3N4 increased, the weight gain of ZrB2-20 vol% SiC composites decreased (from 13.84 to increased (64-128 μm). The 9.84 mg/cm2) and oxide layer thickness weight change of the ZrB2 composites is mainly due to the involvement of different reactions during oxidation (see Table 2). From Table 2, it is clear that the reactions (7), (9), (11) and (12) would be resulting for the increase in weight of ZrB2 composites and reactions (8) and (10) would be leading in weight loss of the samples. The following are the  9069  \\x0c', 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 11. SE micrographs of cross-sectional surface of ZS20-2.5SN after oxidation at 1500 °C for 10 h and its respective layers EDS elemental map.  oxidation reactions (see Table 2):  that are possible at  the oxidation test conditions  ZrB2(s)+5/2O2(g) = ZrO2(s)+B2O3(l)  B2O3(l) = B2O3(g)  SiC(s)+3/2O2(g) = SiO2(l)+CO(g)  SiC(s)+O2(g) = SiO(g)+CO(g)  2ZrN(s)+2O2(g) = 2ZrO2(c)+N2(g)  4BN(s)+3O2(g) = 2B2O3(l)+2N2(g)  (7)  (8)  (9)  (10)  (11)  (12)  From Table 2, except reaction (8) all the other reactions are thermodynamically feasible as they exhibit strong negative free energy. However, thermodynamic models that employ volatility diagrams and kinetic models support evaporation of B2O3 at 1500 °C in air [4]. It is well known that the oxidation of ZrB2 phase lead to the formation of ZrO2 and B2O3. At high temperatures (> 1400 °C), the vaporization of B2O3 takes place and left the porous and non-protective ZrO2 on the surface of ZrB2 [4,26]. In case of ZrB2 composites, the presence of Si containing reinforcements (SiC) helps to form SiO2 and CO. The reaction (9) signifies passive oxidation of SiC and at low oxidation potentials or elevated temperatures the active oxidation of SiC (reaction (10)) takes place. The viscous glassy silica oxide layer shields the surface and limits the inward diffusion of oxygen into the surface of the sample [37-39]. The formation of dense SiO2 layer realized to be beneficial for mechanical properties as well [39]. The other ZrN and BN phases of ZS20-SN composites also takes place in its oxidation (as per reactions (11&12)). For comparison purpose the oxidation characteristics in terms of weight gain, oxide layer thickness of various ZrB2 composites is presented in Table 3 [20,27,40-48]. Talmy et al. studied the effect of addition of SiC, Si3N4, Ta5Si3 and TaSi2 on the oxidation behavior of ZrB2 over a range of temperature 1200-1500 °C [27]. In case of pure ZrB2, low weight loss (6 mg/cm2) and maximum oxide layer thickness (160 μm) was noticed after oxidation at 1300 °C for 2 h (Table 2) [27]. Under similar oxidation conditions, increase in weight gain (14 mg/ cm2) and reduction in oxide layer thickness (140 μm) was noticed when  ZrB2 reinforced with 5 vol% Si3N4. The less weight gain for pure ZrB2 when compared to ZrB2-5vol% Si3N4 was attributed to the more weight loss due to evaporation of B2O3 in pure ZrB2 than the composite. On the other hand, significant increase in weight gain of 25 and 28 mg/cm2 was reported for pure ZrB2 and ZrB2-5vol% Si3N4, respectively when the materials were oxidized at higher temperatures (1500 °C, 2 h). Excellent oxidation resistance or very minimal weight gain (5 mg/cm2) was reported for ZrB2-25 vol% SiC composite [27]. Similarly, very minimal weight gain (5 mg/cm2) was reported for ZrB2-20 vol% SiC-40 vol% MoSi2 even after oxidation at 1500 °C for longer duration of 10 h due to the formation of dense rich SiO2 glassy oxide layers [20]. In a different work, the oxidation behavior of textured and untextured ZrB220 vol% MoSi2 composites was investigated [42,43]. It was reported that the textured ZrB2 composites exhibited better oxidation resistance (weight gain of 7 mg/cm2 after oxidation at 1500 °C for 12 h) than the untextured composites, however, the oxide layer thickness was quite high (155 μm). Moderate weight gain (12.7 mg/cm2) and oxide layer thickness (65 μm) was reported for ZrB2-10SiCf-10ZrSi2 after oxidation at 1500 °C for 10 h [48]. In case of oxidation of pure ZrB2, Fahrenholtz and Hilmas reported formation of small oxide layer thickness of 10 μm thickness of 400 μm (at 900 °C for 8 h) and a very larger oxide layer (1500 °C for 2 h) [49]. By a careful look at Table 3, it can be realized that the oxidation characteristics (weight gain varied from 9.84 to thickness between 64 and 128 μm) of 13.84 mg/cm2 and oxide layer SPSed ZS20-SN composites is comparable with other ZrB2 composites and exhibiting better oxidation resistance. A thorough microstructural investigation of cross-section and surface of the oxidized ZrB2-20SiC-Si3N4 composites were carried out to understand the oxidation behavior. Fig. 10 shows the cross-sectional SEM images of oxidized ZrB2 samples after oxidation at 1500 °C for 10 h. It was observed that all the samples composed of stacking of three different layers, but the thickness of the layers were varying. The ZrB2 composites consist of top silica rich layer, intermediate SiO2-ZrO2 layer and base material (unreacted). From Fig. 10, it can be noticed that the thickness of outer SiO2 layer and intermediate SiO2-ZrO2 layers varying considerably and in particular, its size increased with increasing Si3N4 amount. In the intermediate layer, the ZrO2 content is observed to be  9070  \\x0c', 'S.K. Thimmappa et al.  Ceramics International 45 (2019) 9061-9073  Fig. 12. Surface morphology of (a) ZS20-2.5SN, (b) ZS20-5SN and (c) ZS20-10SN after oxidation at 1500 °C for 10 h, it shows two distinct oxide phases (d) ZrO2 and (e) SiO2.  more with more amount of Si3N4 addition to ZS20 (see Fig. 10 d, f and h). The unreacted beneath layer composed of ZrB2 and SiC (Fig. 10 e, g and i). Fig. 11 shows representative SEM-EDS of cross-sectional microstructure of ZS20-2.5SN after oxidation. The microstructure consists of thick dense outer layer of SiO2, intermediate (ZrO2-SiO2) layer and unreacted bulk. The corresponding EDS spectra also confirmed the presence of ZrO2 and SiO2 phases in the intermediate layer, ZrB2 and SiC phases in the beneath unreacted bulk. In most of the cases, presence of SiC depleted layer along with SiO2 and unreacted base material has been reported for the oxidation of ZrB2-SiC composites [1,2,4,26,49]. The SiC depleted layer lowers the oxidation resistance of ZrB2 composites. However, in the present SPSed ZS20-SN samples, SiC depleted layer was not observed. These observations clearly indicate good oxidation resistance of the ZS20-SN composites. Zhang and Padture have observed improved oxidation resistance for borosilicate glass coated ZrB2-20 vol% SiC composite after oxidation at 1500 °C for up to 20 h duration in air [37]. They also did not report any SiC depleted layer for the composite. In another work, Seong et al. studied the effect of partial pressure of oxygen on SiC (up to 30 vol%) reinforced ZrB2 [38]. Presence of top SiO2 layer, ZrO2-SiO2 mixed layer along with base material without any SiC depletion layer were observed after oxidation at 1500 °C for different time durations (up to 10 h).  The microstructure of oxidized samples surfaces after oxidation (at 1500 °C for 10 h) were shown in Fig. 12. ZrO2 (grey) and SiO2 (dark) phases were noticed in all samples. The presence these phases were confirmed by EDS analysis as well (Fig. 12). The formation of these phases can be understood from the oxidation reactions as presented in Table 2. The oxidation of ZrB2 and ZrN phases results in the formation of ZrO2. From Fig. 12, it is evident the grain size of ZrO2 significantly reduced with increasing Si3N4 content. The grain size of ZrO2 was reduced considerably (from 3.04 μm to 1.91 μm) with the addition of 10 vol% Si3N4 to ZS20. It reflects that the growth of ZrO2 was hindered in ZS20-10SN due to high amount of silica layer, which retards the growth of ZrO2. Zhang et al. [50] reported that ZrO2 particle size increases in ZrB2-SiC-ZrC with increasing oxidation temperature and time. Fig. 13a shows XRD pattern of ZrB2-SiC-Si3N4 composites oxidized surfaces (at 1500 °C). The oxidized ZS20-SN composed of m-ZrO2 and tZrO2 phases. In particular, amount of t-ZrO2 phase increased with increasing the Si3N4 content. The ZS20-10SN composite shows more peak broadening background in the range of 20-30° (diffraction angle), which indicates presence of an amorphous phase. This type of patterns depends on two factors. Firstly, Zr is having a higher atomic number (Z = 40) compared to Si (Z = 14), due to this ZrO2 is having higher atomic scattering which renders to stronger X-ray intensities. The  9071  \\x0c', 'S.K. Thimmappa et al.  Fig. 13. (a) Surface and (b) cross composites at 1500 °C for 10 h.  sectional XRD of oxidized ZrB2-SiC-Si3N4  second factor is that relative amount of ZrO2 at the surface exposed to X-ray is higher in ZS20-2.5SN and ZS20-5SN compared to ZS20-10SN. On the other hand, intensity of zirconia peaks diminishes as the amount of Si3N4 increases, which confirms that amount of silica is high on the surface. The cross-sectional XRD of oxidized samples is presented in Fig. 13b. The presence of ZrB2, SiC and ZrO2 was noticed and the secondary phases were more prominent in ZS20-10SN. Coming to oxidation mechanisms of ZS20-SN, the oxidation of ZrB2 and ZrN phases lead to the formation of porous ZrO2. However, the viscous liquid silica that was formed during oxidation of SiC subsequently fills the pores of ZrO2 and thus presence of continuous layer without any pores and cracks can be observed. In ZS20-10SN, relatively large amount of ZrN present along with ZrB2 major phase when compared to other ZS20 composites. Therefore, more porous ZrO2 expected for ZS20-10SN and these pores subsequently filled by the liquid SiO2. Hence, the thickness of oxide layer was high for ZS20-10SN composite than ZS20-2.5SN and ZS20-5SN. The thickness of silica layer increases with increasing amount of Si3N4 without any spallation/discontinuity, which reduces oxygen diffusion. In fact, the oxygen diffusion co 9072  Ceramics International 45 (2019) 9061-9073  efficient in ZrO2 (10−10 m2/s) is several orders higher in magnitude than SiO2 (10−21 m2/s) at 1550 °C [37]. Hence, SiO2 expected to show more inhibition to the oxygen diffusion when compared to ZrO2. Overall it is noticeable that the ZS20-SN composites exhibited good oxidation resistance as there is no evidence of SiC depleted layer. In the applications point of view, mechanical properties of UHTCs are also need to be considered. As the ZS20-5SN measured with highest hardness, it is expected to be better choice material.  4. Conclusions  (cid:129) The ZrB2-20 vol% SiC-Xvol.% Si3N4 (X = 2.5, 5 and 10 vol%) composites could be densified to more than 98% theoretical density after Spark Plasma Sintering (SPS) at 1900 °C for 3 min under 50 MPa. (cid:129) It has to be noted that even with the use of SPS and Si3N4 a higher sintering temperature of 1900 °C was required to densify ZrB2-20SiC composites. However, the ZrB2-20 vol% SiC-Si3N4 composites were characterized with uniform microstructure and the grain size of ZrB2 varied narrowly between 3.37 and 3.61 μm and it slightly increased with increasing addition of Si3N4. (cid:129) The microstructural analysis of the ZrB2-20 vol% SiC-Si3N4 composites composed of secondary phases (ZrO2, BN, ZrN) along with ZrB2 and SiC major phases. The formation of such new phases indicate involvement of sintering reactions during SPS. All the secondary phases were well dispersed in ZrB2 matrix. (cid:129) The ZrB2-20 vol% SiC-5vol% Si3N4 composite measured with the maximum hardness of 30.56 GPa. In fact, it is the highest ever reported hardness for ZrB2 composites. (cid:129) The weight gain of ZrB2-20 vol% SiC composites decreased (from 13.84 to 9.84 mg/cm2) and oxide layer thickness increased (64-128 μm) with increasing amount of Si3N4 from 2.5 to 10 vol%. It is obvious that the oxidation resistance of ZrB2-20 vol% SiC composites increased with the addition of Si3N4. (cid:129) The microstructure of oxidized samples surfaces after oxidation (at 1500 °C for 10 h) consists of ZrO2 and SiO2 phases in all samples. The cross-sectional microstructure of oxidized ZS20-Si3N4 composites consists of thick dense outer layer of SiO2, intermediate (ZrO2SiO2) layer and unreacted bulk.  Acknowledgement  The financial support by Science and Engineering Research Board (SERB) of Department of Science and Technology (DST), Government of India is gratefully acknowledged. The authors thank TEQIP-II (Centrally Sponsored Scheme (CSS) of MHRD, Govt. of India) for the financial support to procure Thermo-Calc (2017b) software, that was used to plot ternary phase diagram in the present work.  Appendix A. Supplementary data  Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2019.01.243.  References  [2]  [4]  [1] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (5) (2007) 1347-1364. P. Hu, W. Guolin, Z. Wang, Oxidation mechanism and resistance of ZrB2-SiC composites, Corros. Sci. 51 (11) (2009) 2724-2732. [3] D. Sciti, A. Balbo, A. Bellosi, Oxidation behaviour of a pressureless sintered HfB2MoSi2 composite, J. Eur. Ceram. Soc. 29 (9) (2009) 1809-1815. J. 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  "_id": 207,
  "PDF": "Physical characterization and arcjet oxidation of hafnium-based ultra high temperature ceramics fabricated by hot pressing and field-assisted sintering.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 2337-2344  Physical characterization and arcjet oxidation of hafnium-based ultra high temperature ceramics fabricated by hot pressing and field-assisted sintering  Matthew Gasch  , Sylvia Johnson  ∗  NASA Ames Research Center, MS 234-1, Moffett Field, CA 94035, United States  Available online 1 May 2010  Abstract  For this study, HfB2 -based ultra high temperature ceramic (UHTC) samples were prepared by hot pressing and field-assisted sintering (FAS) with 10-20 vol.% SiC (baseline), 5 vol.% TaSi2 , and 5 vol.% iridium. Dense billets were tested for hardness and mechanical strength. When compared, the FAS method consistently yielded materials with a grain size 1.5-2 times finer than samples processed via hot pressing. In general, room temperature flexural strengths of these materials were found to be lower (400 MPa) than similar fully dense HfB2 -SiC materials, with strengths between 500 and 700 MPa. Oxidation resistance testing of flat-face models was conducted in a simulated re-entry environment, at QCold Wall 250 W/cm2 for 5 min. Samples processed by FAS had reduced oxide thickness and SiC depletion zones compared to the baseline HfB2 -20SiC material. In all cases oxide thickness was reduced by 3× and SiC depletion zone thickness was reduced 3× over the baseline. Published by Elsevier Ltd.  Keywords: Hot pressing; Sintering; Hafnium diboride; Oxidation resistance; Composites  1.  Introduction  Ceramic borides, such as hafnium diboride (HfB2 ) and zirconium diboride (ZrB2 ), are members of a family of materials with extremely high melting temperatures referred to as ultra high temperature ceramics (UHTCs). UHTCs constitute a class of promising materials for use in high temperature applications, such as sharp leading edges on future-generation re-entry vehicles, because of their high melting points.1 Despite their potential, wide scale use has been limited, due, in part, to poor fracture toughness and oxidation resistance. Several researchers have looked at modifying the material composition to improve the performance of these materials.2-5,26,27 . However, careful consideration must be given to the grain boundary phases, formed during processing with additives, as these can deteriorate the strength, thermal performance, and physical properties of UHTCs at elevated temperatures.6,7 The controlled development of microstructure has become important to the processing of UHTCs, with the prospect of  ∗  Corresponding author at: NASA Ames Research Center, Thermal Protection  Materials Branch, MS 234-1, Bldg 234, Room 211, Moffett Field, CA 94035,  United States. Tel.: +1 650 604 5377; fax: +1 650 604 0487.  E-mail address: matthew.j.gasch@nasa.gov (M. Gasch).  0955-2219/$ - see front matter. Published by Elsevier Ltd.  doi:10.1016/j.jeurceramsoc.2010.04.019     improving the mechanical and thermal properties of these materials.8-12 The improved oxidation resistance of HfB2 has also become important if this material is to be successfully used at temperatures above 2000 C. The current study has investigated processing HfB2 -based materials with SiC, TaSi2 , and iridium powder. The addition of TaSi2 was pursued in this work based on the previous work of several researchers. Talmy et al. investigated the oxidation of ZrB2 ceramics with SiC, Si3N4 , Ta5Si3 and TaSi2 additions and found that all additions improved the oxidation resistance of ZrB2 /SiC ceramics below 1400 C.19 They found that improved oxidation resistance correlated with modification of the chemical composition of the surface oxide layer, leading to decreased inward diffusion of oxygen. Furthermore Peng and Speyer demonstrated that TaSi2 additions improved the oxidation of ZrB2 materials up to 1550 C.18 Iridium is also well known for its high temperature oxidation resistance, moreover the addition of metallic or semi-metallic sintering aids such as Fe, Ni, Co, W, WC have been shown to improve final density of Hf and Zr-based ceramics and these additions allow for lower densification temperatures.21 In addition to the investigation of additives for improving oxidation resistance and processing temperature of Hf-based composites, two processing techniques were evaluated in this work; conventional hot pressing as well as electric field-assisted        \\x0c', '2338  M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  sintering (FAS). The resulting microstructural variations were evaluated with SEM; the mechanical strength of each sample was evaluated using bi-axial flexure; and oxidation behavior of these materials was evaluated in a simulated re-entry environment in the AHF arcjet facility at NASA Ames Research Center, during a 5-min exposure to a cold wall heat flux of 250 W/cm2 and stagnation pressure of 0.1 atm.  2. Experimental procedure        The raw powders used in this work were: −325 mesh HfB2 Inc., 1-2 \\u242em SiC from H.C. Starck Inc., and −325 mesh iridium from SurePure. Analysis of and TaSi2 from Cerac the crystalline phases present in the raw and mixed powders was performed with X-ray diffraction (Scintag X-ray diffractometer) using Cu k␣ radiation. Particle size of the raw and milled powders was measured using laser light scattering (Malvern). Powders were then weighed according to volume % of each material and then wet-milled them with WC(Co) milling media in a planetary mill (Fritsch Pulverisette 5, Germany). The milled powders were carefully dried to prevent phase segregation between the HfB2 (ρ = 11.12 g/cm3 ) and the SiC (ρ = 3.2 g/cm3 ). After drying, the powders were loaded into a 25 mm diameter graphite die lined with graphfoil. Hot pressing was conducted in a graphite element resistance furnace (Thermal Technologies HP50-7010G, Santa Rosa, CA) with vacuum levels of <200 mTorr. Above 1600 C, the partial pressure of carbon in the furnace increased significantly and degraded the vacuum. Consequently, we back-filled the furnace with one atmosphere of inert gas (argon or helium), to preserve the graphite element and insulation. Typical furnace conditions required for densification were 1900-2200 C for 1 h at 25 MPa. Field-assisted sintering (FAS) was conducted at the AFOSR labs in Dayton, OH. The FAS system there is a FCT System GmbH Model HPD 25-1 (Rauenstein, Germany). We used the same milled powders as those densified by conventional HP. Samples were again loaded into a 25 mm diameter graphite die lined with graphfoil. Typical conditions required for FAS densification were 1800-1900 C, with hold times of 5-10 min. Ramp rates used during FAS ranged from 100 to 300 C per eter focused on the bottom of a bore hole in the punch 5 mm minute. Sample temperature was measured by an optical pyromfrom the powder. Density of the hot-pressed billets and test specimens was measured using the Archimedes method. After the samples were cross-sectioned and polished to a 1-\\u242em finish, we characterized the microstructure of consolidated samples using optical microscopy, scanning electron microscopy (FEI ESEM 30, Hillsborough, OR) with EDX analysis and transmission electron microscopy (FEI TEM CM200FEG, Philips, Eindhoven, Netherlands) operating at 200 kV. Average grain size was determined using the lineal intercept method without grain shape corrections, following ASTM E112. TEM (FEI TEM CM200FEG, Philips, Eindhoven, Netherlands) specimens were prepared using a method developed for thinning ceramic fibers, but the technique is equally applito bulk samples.22,23 Samples were cable sectioned with a           low-speed diamond saw into thin slices and then mechanically ground and polished to 5-10 \\u242em thickness with the aid of a tripod polisher using successively finer diamond lapping films. Final thinning was accomplished by mounting the thin sections on 50-mesh copper grids followed by low-angle ion milling (Precision Ion Polishing System, Gatan, Pleasanton, CA) to electron transparency. Energy-dispersive X-ray spectroscopy (Model Voyager, Noran Instruments, Middleton, WI) was used for elemental analysis. Mechanical and thermal property test specimens were prepared with diamond tooling and ground to a final surface finish in accordance with ASTM C1161 (Chand Kare Technical Associates, Worcester, MA). Vickers hardness was determined using indentations created with a Shimadzu indenter (HSV-30, Japan). Hardness values were calculated from the average length of the diagonal lines across an indentation made by applying a load of 49 N over a 15-s interval. Flexural strength was measured using bi-axial flexure testing, according to ASTM 1499 “Monotonic Equibiaxial Flexure Testing,” using 25.4 mm diameter × 2 mm thick disks with a crosshead speed of 0.5 mm/min. The laser flash diffusivity technique was used (Netzsch LFA457, Germany), over the range 25-600 C to measure thermal diffusivities. Samples were coated with a light spray of carbon powder to prevent reflection of the incident laser beam. Measurements were made under vacuum, using a laser (1538 V setting) with a pulse duration of 0.5 ms. Signals were corrected to account for heat losses using Cowen + Pulse correction method.25 We repeated thermal diffusivity measurements three times at each temperature, and averaged the results, arriving at extremely repeatable measurements, with typical standard deviations of less than 0.5%. Conversion to thermal conductivity requires numerical values for the density and specific heat at the corresponding temperatures. We used the measured room temperature densities of each sample, considering the effect of thermal expansion range insignificant over the temperature range measured. The method used to calculate specific heat values for all samples is reported elsewhere.24 Finally, oxidation resistance was conducted on machined flatface models in a simulated re-entry environment created via the NASA Ames Advanced Heating Facility (AHF) arcjet. A photo of an as-machined flat-face arcjet model is shown in Fig. 1. The model is 25.4 mm in diameter, and the overall height is 8 mm. The notch in the base of the model, shown in the lower left of Fig. 1, is used to pin the model into the holder. Models were placed in SiC-coated graphite holders (shown in the right-most image of the same figure), enabling test durations in excess of 10 min. Models and holders were then attached to a water-cooled arm (sting) that moved the models in and out of the plasma stream. A variety of instrumentation was used to calibrate arcjet conditions and measure the thermal response of the materials. Cold wall heat flux values, shown in Table 1, were measured using a copper Gardon gauge and are referenced to a 76 mm diameter hemisphere. However, hot wall heat flux values at the model’s surface differ, because of differences in model geometry and between the catalycity of the models and the catalycity of the copper Gardon gauge (detailed discussion of the heat flux at the  \\x0c', 'M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  2339  Fig. 1. Flat-face arcjet model shown as fabricated, attached to SiC-coated graphite holder and attached to water-cooled sting arm.  surface of the model is beyond the scope of this work). Twocolor optical pyrometers were used to make surface temperature measurements during the tests. Oxide thickness of post-arcjet models was determined after models were cross-sectioned and polished to a 1-\\u242em finish and then inspected under SEM as described above. Oxide thickness was determined by averaging 5-8 thickness measurements across an image that was taken at the middle of the model (12 mm from either edge).  3. Results and discussion  3.1. Processing and properties  Table 1 shows the average particle size, crystalline phases present, and surface area of the as-received powders. After milling, the average particle size of the powder mixtures is 2.1 \\u242em. Table 2 lists the physical characteristics of the consolidated samples. The densities of consolidated samples ranged from 94 to 99% of the theoretical density (TD). Fig. 2 shows SEM micrographs of polished cross-sections. The dark phase in the images is SiC and the lighter phase is HfB2 . TaSi2 additions are hard to distinguish from the HfB2 phase, because of the similar atomic weight; however, for the sample with iridium, the iridium phase is the brightest color and is clearly distributed at grain boundaries and triple points in the hotpressed sample. The FAS sample with iridium had much less time for solid-state diffusion, and the iridium grains appear more equiaxed, like that of the hafnium phase. Upon analysis of the billet post-processing, cracks were visible in the FAS sample with iridium, within the bulk material; thus, we did not mea Table 1  Summary of the properties of the starting powders used in this study.  such as hardness and strength, on this  sure some properties, material. Table 2 summarizes the experimental values of some mechanical properties. Vickers hardness values for the materials processed during this work are between 13 and 18 GPa. The hot-pressed samples with TaSi2 and iridium have hardness values on the low end, in part due to the slight porosity of those samples. However, the FAS sample with TaSi2 had the highest hardness, possibly a result of a reduced grain size, despite also having the lowest theoretical density. Average strengths of the baseline HfB2-SiC material ranged from 428 to 485 MPa, for the FAS and hot-pressed samples, respectively. Average strengths for the samples with TaSi2 ranged from 330 to 380 MPa; again the hot-pressed material had the highest strength. It is interesting to note that the addition of TaSi2 to the baseline material resulted in lower average strengths. The addition of 5 vol.% iridium and TaSi2 to the baseline resulted in hot-pressed samples with a reduced grain size. However, the same materials processed via FAS cracked; thus, hardness and strength were not measured. In general, the room temperature flexural strengths of these materials is rather low compared to results reported on similar fully dense HfB2 -SiC materials.9,17 In fact, samples processed by FAS consistently had lower strength values, despite having refined grain size. These results are somewhat contrary to the expected result; higher strengths were expected for the samples processed by FAS. Research by Guo et al.12 on ZrB2 -based composites found that the flexural strength of samples fabricated with nano-sized SiC particles was much higher than samples processed with micronsized SiC. The observed increase in strength was attributed to SiC particles within ZrB2 grains forming in situ inter-granular  Vendor  Starck  Starck  Cerac  SurePure  Powder  HfB2 SiC  TaSi2 Ir  Surface area (m2 /g)  Particle size (d0.5) (\\u242em)  Crystalline phases detected  0.43  8.09  -  0.15  4.1  1.6  4.6  11.6  HfB2 /HfC SiC  TaSi2 Ir  \\x0c', '2340  Table 2  M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  Summary of the physical characterizations done on consolidated samples processed in this study.  Sample ID  Hot-pressed properties  HfB2 -20SiC (Baseline) HfB2 -10SiC-5TaSi2 HfB2 -15SiC-5TaSi2 -5Ir  Density (g/cm3 )  9.49 (99% TD)  9.45 (96% TD)  9.98 (96% TD)  Grain size (\\u242em)  Hardness (HV)  Strength (MPa)  7.7  8.5  5.1  16.5  13.0  13.0  485  380  -  Sample ID  Field assist sintered properties  HfB2 -20SiC (Baseline) HfB2 -10SiC-5TaSi2 HfB2 -15SiC-5TaSi2 -5Ir  Density (g/cm3 )  9.46 (99% TD)  9.61 (94% TD)  10.02 (97% TD)  Grain size (\\u242em)  Hardness (HV)  Strength (MPa)  4.1  2.3  1.6  17.5  18.0  -  428  330  -  SiC, TaSi2 and Ir amounts represent volume % of each addition.  composites. Thus, while the fabrication of UHTC composites by FAS creates refined grains, an improvement in flexural strength is likely not possible without the fabrication of materials using nano-sized starting powders. Thermal conductivity measurements were performed on all materials, as a function of temperature from 25 to 600 C; see Fig. 3. There is a significant difference between the pure HfB2 , fabricated in a previous study from elemental powders via FAS,27 and the baseline HfB2-SiC materials fabricated at ARC. Materials originally hot pressed at ARC under a previhas a room temperature conductivity of 40 W/mK. ous program used a long process that resulted in material that In this study, in an 2× increase in conductivity, the processing of the baseline material via FAS results to 80 W/mK at room     temperature. Preliminary research indicates that the migration of impurities from grain boundaries into the bulk HfB2 during long hot-pressing times are the main reason for reduced thermal conductivity.24 As shown in Fig. 4, TEM analysis demonstrates that Mg, Al and oxygen contaminates could be found in several locations. Cobalt contamination (from the WC(Co) milling media) was also detected at grain boundaries, however, tungsten and tantalum were not found in grain boundaries. Increased processing times could allow the electronic structure/vacancy concentration of bulk HfB2 and SiC to be altered either via the diffusion of impurities from grain boundaries into the bulk or by the loss of boron or carbon during processing.24,27 The baseline HfB2-SiC materials presented in this paper were processed with a modified hot-press schedule that, while shorter than the  Fig. 2. SEM images of the current UHTC samples comparing the differences in microstructure between hot pressing and field-assisted sintering.  \\x0c', 'M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  2341           sample can be attributed to the formation of a more uniform oxide scale, whereas the FAS sample appears to have oxidized less across the surface. Post-test emittance measurements confirm that the FAS sample has a thinner oxide layer, as this sample had a much higher emittance. A similar trend was observed for the samples with TaSi2 additions, where the surface temperature for the hot-pressed sample was 45 C higher than the FAS sample. The sample with TaSi2 and Ir additions exhibited surface temperatures 100 C lower than the baseline hot-pressed material and 30 C higher than the sample with TaSi2 additions and no Ir additions. Interestingly, the surfaces of samples with SiC and TaSi2 additives remained smooth during testing and after cooldown. However, droplets were observed to have formed on the sample with Ir during testing. In the post-test surface image, droplets of iridium (confirmed by EDX analysis) dot the otherwise smooth surface of the model. In previous studies by ManLabs and others, the oxidation of HfB2 /SiC samples was observed to leave three distinct regions within the material.1,13-16 The first region comprises the surface oxide, primarily composed of SiO2 and some HfO2 . Below that is a SiC-depleted zone, where the SiC has oxidized away (active oxidation), leaving behind a porous HfB2 matrix. Below the depletion zone is virgin material. The flat-face models tested in this series were cross-sectioned to determine if a similar structure had formed. Fig. 6 shows the SEM images of the cross-sections of each arcjet sample. Each sample clearly shows a white SiO2 oxide layer, below which is a SiC-depleted region of porous HfB2 grains. A comparison of the measured oxide thickness and SiC depletion zone thickness for each sample is listed in Table 3. In general, the oxidation resistance of samples with TaSi2 additions is better than the baseline HfB2-SiC materials, as expected.18,19 A comparison of hot-pressed materials shows the addition of TaSi2 reduced  Fig. 3. Thermal conductivity as a function of temperature for each of the mate rials presented in this paper.  previous schedule, is 6× longer than samples processed by FAS. Interestingly, baseline (HfB2-20SiC) materials hot pressed with the shorter schedule demonstrate thermal conductivity similar to materials processed by FAS. However, the addition of TaSi2 and/or iridium to the baseline powder yields material with a conductivity reduced to 50 W/mK at room temperature.  3.2. Arcjet testing  Fig. 5 provides a comparison of the surfaces of each of the flat-face models after arcjet exposure at 250-280 W/cm2 for 5 min. A comparison of the measured surface temperatures and post-test emittance is listed in Table 3. For the baseline HfB2 /SiC samples, steady state surface temperatures were 1690 the hot-pressed material and 1530 C for C for the FAS processed sample. The higher temperature observed on the hot-pressed        Fig. 4. TEM images of the current and spectra of some impurities at the grain boundaries of hot-pressed HfB2 -20SiC material processed using a long hold time.  Table 3  Summary of the arcjet conditions, surface temperature, post-test emittance and oxide thickness of cross-sectioned arcjet models.  Sample ID  HfB2 -20SiC (Baseline) HfB2 -20SiC (Baseline) HfB2 -10SiC-5TaSi2 HfB2 -10SiC-5TaSi2 HfB2 -15SiC-5TaSi2 -5Ir  Sintering  method  Hot press  FAS  Hot press  FAS  Hot press  CW heat flux (W/cm2 )  280  250  250  250  250  Pstag (atm)  Test duration (s)  Surf. temp. (     C)  Post-test  emittance  Oxide thickness (\\u242em)  SiC depletion (\\u242em)  0.19  0.10  0.10  0.10  0.10  600  600  600  600  600  1690  1530  1560  1515  1590  0.67  0.87  0.89  0.89  0.87  13  3  7  3  4  24  8  34  6  9  \\x0c', '2342  M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  Fig. 5.  Images of the surface of each of flat-face models after 5-min arcjet exposure.  oxide thickness by 2× but, the SiC depletion zone was 1.5× larger than the baseline material. Interestingly, the addition of TaSi2 and Ir reduced oxide thickness by 3×. The iridium appears insoluble in HfB2 and can be observed along grain boundaries and triple points, perhaps acting as an oxygen barrier around HfB2 grains. All samples processed by FAS had reduced oxide thickness and SiC depletion zones. In all cases, oxide thickness was reduced by 3× and SiC depletion zone thickness was reduced 3× over the baseline. Very recently, Hwang et al.20 reported that the incorporation of nano-sized SiC particles improved the oxidation resistance of hot-pressed ZrB2 ceramics. The results of our work appear to support their hypothesis, that ceramics  with reduced grain size have an increased diboride/SiC interface length per unit area of exposed surface and a decreased spacing between Si-containing particles. The decrease in spacing between SiC particles allows the surface to more rapidly form a protective oxide over the diboride phase. More arcjet testing is required to improve our understanding of the oxidation mechanisms within these materials in high temperature oxidizing environments. Additional testing is also required to verify the reproducibility of the current results. However, these results, though preliminary, suggest that further improvements in the oxidation resistance of HfB2 materials can be realized by achieving small-grained microstructures with well distributed SiC or TaSi2 dispersions. Finally, the addition  \\x0c', 'M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  2343  Fig. 6. Cross-sectional images of each of the arcjet models post-test.  of iridium also appears to reduce oxidation of these materials perhaps by acting as an oxygen barrier around HfB2 grains.  4. Conclusions  Three different ultra high temperature ceramics: HfB2 + 20 vol.% SiC (baseline), HfB2 + 10 vol.%SiC + 5 vol.%TaSi2 , and HfB2 + 5 vol.% TaSi2 + 5 vol.% iridium, were prepared by hot pressing and field-assisted sintering sistently yielded materials with a grain size 1.5-2× finer than (FAS). In comparison to hot pressing, the FAS method consamples processed via hot pressing. Consolidated specimens were tested for hardness and mechanical strength via bi-axial 25 mm diameter × 2 mm thick flexure of disks. Samples processed by FAS consistently had lower strength values, despite having refined grain size, and, in general, the room to be low (400 MPa) compared to results reported on similar temperature flexural strengths of these materials were found fully dense HfB2-SiC materials that showed strengths between 500 and 700 MPa. Oxidation resistance was characterized in a simulated reof flat-face models was conducted at QCold Wall 250 W/cm2 entry environment in the Ames AHF arcjet facility. Testing for 5 min. All samples processed by FAS had reduced oxide thickness and thinner SiC depletion zones than the baseline HfB2 -20SiC material reduced by 3× (from 13 to 3 \\u242em) that was hot pressed. In all cases, oxide SiC depletion zone thickness was reduced 3× (from 24 to thickness was and 8 \\u242em) over the baseline. The results of the work presented  here appear to support the work of others, whereby diboride ceramics fabricated with reduced SiC grains have an increased diboride/SiC interface length per unit area of exposed surface and a decreased spacing between Si-containing particles. The decrease in spacing between SiC particles allows the surface to more rapidly form a protective oxide over the diboride phase.20  Acknowledgments  This work was funded by the Hypersonics element of the Fundamental Aeronautics Program at NASA Aeronautics Research Mission Directorate (ARMD). Michael Cinibulk of the Materials and Manufacturing Directorate, Air Force Research Laboratory, performed the TEM work. We also acknowledge NASA-SCAP for their critical financial support of the arcjet operational capability at Ames. Portions of this work were also performed under NASA contract NAS2-99092 to ELORET Corp. We would like to thank Jerry Ridge of ELORET for his support measuring the surface emittance of UHTC arcjet samples.  References  1. Kaufman L, Clougherty EV.  Investigation  of  boride  compounds  for  high temperature applications, RTD-TRD-N69-73497. Part XXXVII. Cam bridge, MA: ManLabs Inc.; 1963.  2. Guo S-Q. Densification of ZrB2 -based composites and their mechanical and physical properties: a review. J Eur Ceram Soc 2009;29:995-1001.  3. Guo W-M, Vleugels J, Zhang G-J, Wang P-L, Van der Biest O. Effects  of Re2O3 (Re = La, Nd, Y and Yb) additions in hot-pressed ZrB2-SiC ceramics. J Eur Ceram Soc 2009;29:3063-8.  \\x0c', '2344  M. Gasch, S. Johnson / Journal of the European Ceramic Society 30 (2010) 2337-2344  4. Li G, Han W, Zhang X, Han J, Meng S. Ablation resistance of ZrB2 -Sic-AlN ceramic composites. J Alloys Compd 2009;479:299-302. 5. Sciti D, Silvestroni L, Nygren M. Spark plasma sintering of Zrand Hf borides with decreasing amounts of MoSi2 as sintering aid. J Eur Ceram Soc 2008;28:1287-96. 6. Monteverde F, Bellosi A. J Mater Res 2004;19:3576. 7. Monteverde F, Guicciardi S, Bellosi A. Mater Sci Eng A 2003;34:6310.  8. Ni D-W, Zhang G-J, Kan Y-M, Sakka Y. Textured HfB2 -based ultra high temperature ceramics with anisotropic oxidation behavior. Scripta Mater 2009;60:913-6.  9. Monteverde F. Progress  in  the  fabrication  of  ultra  high  temperature  ceramics: ‘in-situ’ synthesis, microstructure, and properties of a reactive hot-pressed HfB2 -SiC composite. Compos Sci Technol 2005;65:1869-79. 10. Chen L, Gu Y, Shi L, Yang Z, Ma J, Qian J. Synthesis and oxidation of nanocrystalline HfB2 . J Alloys Compd 2004;36:8353-6. 11. Zhang X, Xu L, Du S, Han W, Han J. Crack-healing behavior of zirconium  diboride composite reinforced with silicon carbide whiskers. Scripta Mater 2008;59:1222-5.  12. Guo S-Q, Yang J-M, Tanaka D, Kagawa Y. Effect of thermal exposure on  strength of ZrB2 -based composites with nano-sized SiC particles. Compos Sci Technol 2008;68:3033-40.  13. Gasch M, Ellerby D, Irby E, Beckman S, Gusman M, Johnson S. Process ing, properties and arcjet oxidation of hafnium diboride-silicon carbide ultra high temperature ceramics. J Mater Sci 2004;39:5925-37.  14. Monteverde  F.  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Ablation behavior of ZrB2 -SiC ultra high temperature ceramics under simulated atmospheric Re-entry conditions. Compos Sci Technol 2008;68:1718-26.  ultra high temperature 2004;39(19):5969-77.  ceramics:  effect  of ZrB2 and HfB2 -based of Ta additions.  J Mater  Sci  27. Zhang SC, Hilmas GE, Fahrenholtz WG.  Improved oxidation resistance  of zirconium diboride by tungsten carbide additions. J Am Ceram Soc 2008;91(1):3530-5.  \\x0c']"
},{
  "_id": 208,
  "PDF": "Plasma wind tunnel testing of ultra-high temperature ZrB2–SiC composites under hypersonic re-entry conditions.pdf",
  "Text": "['Available online at www.sciencedirect.com  Journal of the European Ceramic Society 30 (2010) 2313-2321  Plasma wind tunnel testing of ultra-high temperature ZrB2-SiC composites under hypersonic re-entry conditions  Frederic Monteverde a,∗  , Raffaele Savino b , Mario De Stefano Fumo b , Andrea Di Maso b  a Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy b Department of Aerospace Engineering, University of Naples “Federico II”, P.le Tecchio 80, 80125 Naples, Italy  Available online 20 February 2010  Abstract  The resistance to oxidation and optical properties of a hot-pressed ZrB2 -SiC composite were studied under aero-thermal heating in a strongly dissociated flow that simulates hypersonic re-entry conditions. Ultra-high temperature ceramic models with a blunt or sharp profile were exposed to high enthalpy flows of an N2 /O2 gas mixture up to 10 MJ/kg for a full duration of 540 s, the surface temperatures approaching 2100 K. Stagnation-point temperatures as well as spectral emissivities were directly determined using an optical pyrometer. Microstructural features of the oxidized layers were correlated to optical properties through computational fluid dynamics calculations which allow for numerical rebuilding of key parameters like surface temperatures, wall heat fluxes, shear stresses or concentrations of the species composing the reacting gas mixture. Gradients of temperature on the surfaces facing the hot gas flow established different boundary conditions that led to the formation and evolution of distinct layered oxide scales. © 2010 Elsevier Ltd. All rights reserved.  Keywords: Hot-pressing; Borides; Electron microscopy; Thermal properties; Arc-jet testing  1.  Introduction  Ultra-high temperature ceramics (UHTCs) containing transition metal borides are under intensive analysis for applications in extreme environments,1-3 where high-temperature aerospace applications on hypersonic vehicles have been the most investigated. UHTC composites based on ZrB2 and HfB2 , in combination with silica formers such as SiC or metal silicides like MoSi2 , TaSi2 , or HfSi2 , are particularly attractive for leading edge and control surface components on hypersonic vehicles.4-6 To reduce aerodynamic drag and thus enhance lift-to-drag and maneuverability, hypersonic vehicles require sharp leading edged profiles with radii of curvature on the order of a few millimeters. This configuration conversely gives rise to harsher heat fluxes at the stagnation point, on the order of some MW/m2 , that in turn lead to generated surface temperatures exceeding the usage temperature of conventional aerospace materials, such reinforced SiC matrix composites.7 The utias carbon fiber lization of materials with high thermal conductivities at high  temperatures, like ZrB2 and HfB2 , is an option that favors a faster spatial redistribution of the excess heat, and therefore may ensure performance advantages for sharp leading-edge components for which drawing heat away from the stagnation points is essential. Lower manufacturing costs and reduced density make ZrB2 -based composites more attractive for aerospace applications, compared to their HfB2 -based analogs. In spite of the high melting points for ZrB2 and its metal oxide ZrO2 , both above 2800 K, the technological solution is still challenged by the adverse consequences of oxidation and levels of fracture toughness unable to prevent cracking during rapid thermal transients. The introduction of SiC in combination with ZrB2 has demonstrated enhanced resistance to oxidation, fracture toughness and thermal shock resistance.8,9 A large body of research has also been addressing the oxidation behavior of ZrB2 and its composites using conventional air-furnace environments.10-13 In high-temperature oxidizing environments, ZrB2 -SiC composites react with oxygen through the net parallel reactions: ZrB2 + 5/2O2 = ZrO2 + B2O3 SiC + 3/2O2 = SiO2 + CO(g)  (1)  (2)  ∗  Corresponding author. Tel.: +39 0546 699758; fax: +39 0546 46381.  E-mail address: frederic.monteverde@istec.cnr.it (F. Monteverde).  that yield oxide by-products like zirconia (ZrO2 ), boron oxide (B2O3 ) and silica (SiO2 ). Boron oxide, which melts at 450 C,     0955-2219/$ - see front matter © 2010 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2010.01.029  \\x0c', '2314  F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321     promptly mixes with SiO2 for temperatures above 1200 C to form a borosilicate glass that spreads over and seals the external surfaces. The large volume increase upon oxidation of the bulk material due to the formation of solid zirconia was proposed as the driving force for the upward transportation of the fluid borosilicate glass.10 With increasing testing temperature, the B2O3 component of the borosilicate glass possesses an unusually high vapour pressure that leads it to evaporate preferentially from the glassy phase leaving behind a boron-deficient silica layer. As the silica-rich layer thickens, it slows down the inward diffusion of oxygen to the pristine material below, limiting oxidation and lowering the oxygen partial pressure in the reaction zone.14 At sufficiently low oxygen pressures,15 the oxidation of ZrB2 becomes negligible and the oxidation of silicon carbide becomes “active,” proceeding through the main reaction:  SiC + O2 (g) = SiO(g) + CO(g)  (3)  In air-furnace studies of ZrB2-SiC composites typically have ranged from 1273 to 1873 K, and more rarely up to 2173 K,12,13 at atmospheric pressures. When oxidized, this group of UHTC ZrB2 -composites (10-30% SiC) forms a multilayer oxide scale with an outermost dense silica-rich glassy layer decorated with ZrO2 crystallites and an inner SiC-depleted diboride layer, separated at times by an additional layer based on ZrO2 crystals embedded in a silica-rich glassy network. Test conditions such as temperature and time, or type of secondary phases (coupled to initial sintering additives) affect compositions and relative thicknesses of the reaction layers.16-18 Oxidation rates, most often quantified by monitoring the mass change of the specimens in a thermal gravimetric analyzer working at atmospheric pressure, are consistent with diffusion-limited oxidation (i.e., parabolic kinetics). Oxidation under lower total pressures and higher temperatures can result in evaporation of the SiO2 in addition to the B2O3 , leaving behind a single outer layer of porous ZrO2 and resulting in rapid, linear oxidation kinetics.15 In contrast to conventional high-temperature air-furnace experiments where the gas environment can be assumed in thermal equilibrium with the test specimen, arc-jet wind tunnels generate high-enthalpy and strongly dissociated gas flows that more actively reproduce the aero-heating expected in service on hypersonic vehicles. Such an aero-thermal environment involves not only large temperature gradients between the specimen surfaces and the boundary layer and shock layer edges, but also an energized reactive gas mixture composed of ions, atoms, and molecules in highly energetic states whose vibrational temperature may exceed 10,000 K. The shocked gas will undergo thermochemical relaxation as it nears the colder surface, but typically will not reach chemical equilibrium at the temperature of the surface. One implication of this thermochemical non-equilibrium is that UHTC materials on leading-edge surfaces may interact with significant concentrations of dissociated gas species, whose reactivities often trigger new exothermic reactions. The validation of this class of materials to operate in such sustained environments requires a test campaign in ground-simulated (hypersonic) re-entry conditions, using facilities capable of simulating at least the representative total  enthalpy, pressure and heat flux, even though they cannot simultaneously simulate the fully representative flight envelope in terms of Mach and Reynolds numbers.19 A restricted number of available facilities in conjunction with expensive testing costs still limits the execution of dedicated studies. In this paper, the oxidation behavior of a hot-pressed ZrB2-SiC composite shaped as a blunt hemisphere or sharp cone was investigated using a plasma wind tunnel facility which simulates hypersonic re-entry conditions. The microstructure and composition of the resulting oxide layers were characterized and discussed. Surface temperatures and spectral emissivities were directly measured, and input in computational fluid dynamics (CFD) simulations to rebuild the aero-heating environment.  2. Materials  The ZrB2 + 20 vol% SiC UHTC composite was prepared from commercially available powders supplied by H.C. Starck (ZrB2 grade B, ␤-SiC BF12). An amount of 3 vol% Si3N4 supplied by Bayer (type Baysinid) was added into the base mixture as a sintering aid. Powders in due proportions were ball-milled for 24 h in a polyethylene bottle using silicon nitride media and absolute ethyl alcohol, dried by a rotary evaporator, and sieved (250 \\u242em mesh size). The dried powder was cold compacted into a pellet (44 mm diameter) using an uniaxial press and then positioned inside a graphite mould whose internal walls were lined with BN-coated graphite foil. Hot-pressing of the powder pellet was carried out under rough vacuum (0.5-1 mbar) heating up to 2173 K (about 20 K/min heating rate and 30 MPa of applied pressure). The isothermal hold at 2173 K lasted about 10 min. At the end of the hold, the applied pressure was released and the hot press was cooled to room temperature. After removal from the dies, the hot-pressed compact was about 20 mm thick. The ceramic samples for the aero-thermodynamic study were first cut out by electrical discharge machining and then finished down to the desired dimensions using diamond-loaded tools. The base diameter is 10 mm for both the specimens, whilst the full lengths (specimen plus afterbody) are 10 mm (hemisphere) and 11 mm (cone). The curvature radii of the hemisphere and cone are, respectively, 5 and 0.5 mm.  3. Experiment  3.1. Plasma wind tunnel gallery and test conditions  The experiments were performed in the arc-jet facility “Small Planetary Entry Simulator” available at the Department of Aerospace Engineering, University of Naples “Federico II”. The facility is equipped with a 80 kW plasma torch that operates with inert gases (He, N2 , Ar and their mixtures) at mass flow rates up to 5 g/s. In this case only nitrogen (80 wt%) and oxygen (20 wt%) were used. All the experiments have been carried out with a nitrogen mass flow rate of 0.8 g/s and oxygen mass flow rate of 0.2 g/s. To replicate the air composition, a percentage of 20 wt% molecular oxygen is mixed to the primary nitrogen jet into a mixing chamber after the torch. The nozzle has a throat diameter of 11 mm and an outer diameter of 22 mm. The nominal  \\x0c', 'F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  2315  Mach number (M), determined by the nozzle area ratio, is around 3. The ceramic model is positioned at a distance of 10 mm from the nozzle exit. Pressure transducers are installed at the nozzle exit and inside the test chamber, whilst the stagnation-point pressure at the model location is measured by a dedicated probe. The arc-power and cooling power are controlled separately, and an integrated electronic system also controls gas feed and torch current. The specific total enthalpies H0 at the torch and nozzle exit are calculated through an energy balance on the system torchnozzle, knowing the arc power and the cooling power. It must be pointed out that the measured total enthalpy is an average value for the flow (sometimes referred to as bulk enthalpy), and also that arc-jet facilities typically have radial gradients of the thermo-fluid-dynamic parameters, for instance the enthalpy.20 The test chamber is equipped with portholes and windows for optical diagnostics. Sample surface temperature is measured by a two-colour pyrometer (Infrather ISQ5, Impac Electronic GmbH, Germany) at an acquisition rate of 100 Hz. The pyrometer uses two overlapping infrared wavelength bands at 0.8-1.1 and 1-1.1 \\u242em and measures temperatures from 1300 to 3300 K with a 1% measurement accuracy in the range 1720-3300 K. The ISQ5 pyrometer can be used both in two-colour and one-colour modes. The two-colour mode is used to measure the surface temperature, whilst the single colour mode allows to determine the emissivity once the temperature is known. The arc-jet facility presently available has no optical access in front of the model and therefore the surface model is pointed from the side of the test chamber. The measuring spot of the pyrometer is 3 mm wide and provides an average temperature of this area. For the hemispheric and conical samples the pyrometer pointed, respectively, at 50% and 80% of the radius distances from the tip. The arc-jet wind tunnel, when used in the supersonic flow regimes, feeds a non-equilibrium flow to the shock edge (i.e. at the model’s location): this might have introduced some uncertainties in the estimates of the actual gas chemical composition at the specimen surface.  3.2. Computational fluid dynamics (CFD) and thermal model  Sample temperature, gas pressure, and chemical composition of the gas mixture interacting with the UHTC model surface are key environmental parameters that influence oxidation. Surface temperature and gas pressure are determined directly, whilst the gas composition at the sample surface is computed through the numerical model based on the solution of the Favre-averaged Navier-Stokes equations for a mixture of reacting species in chemical and vibrational non-equilibrium.21 Species viscosity, thermal conductivity, and mass diffusivity are derived from the kinetic theory of gases22 as functions of the Lennard-Jones parameters. The solver computes the mixture viscosity and thermal conductivity with the semi-empirical Wilke’s rule. The flow is considered as a mixture of five gas species: N, O, N2 , O2 , and NO. From the thermodynamic viewpoint, the system is considered a mixture of reacting ideal gases. In the present computations, the molecular species have translational and rotational modes in equilibrium, but the vibrational energy can be at dif ferent temperatures (TV ). For the chemical reactions, the Park model23 is included, with the reaction rate constants specified by the Arrhenius law. For the vibrational and thermal relaxations, the Landau-Teller model was used,24 with a Park correction for the higher temperatures.25 At solid walls, the no-slip condition is enforced by imposing the velocity components to zero. In chemically reacting flows, the mass fractions of the species are dependent variables with their transport equation. Species boundary conditions at the wall are assigned according to the behavior of the solid surface. For a fully catalytic wall, the chemical reactions are catalyzed at an infinite rate and the mass transport at the wall is limited only by diffusion, therefore the mass fractions of the dissociated species at the wall are set equal to zero. For a non-catalytic wall, the diffusive flux of atoms at the wall is set to zero. In addition, a user-defined function has been developed to simulate a wall with a finite value of surface catalycity,26 according to Eq. (4):  ∂Ci  ∂n  = CiKw Di  (4)  where n is the normal coordinate to the wall, Ci is the atomic mass fraction, Di is the species diffusion coefficient and Kw is of the wall material recombination coefficient γ w (0 ≥ γ w ≤ 1), the material catalytic constant. The Kw parameter is a function of the chemical species and of the temperature, according to Eq. (5):  (cid:2)  Kw = γw  R0Tw  2πmi  (5)  R0 is the universal gas constant, mi is the atomic mass. In the present work the same recombination coefficient has been considered both for Oxygen and Nitrogen. The vibrational energy at the wall is set considering vibration thermodynamic equilibrium. Convective fluxes were computed according to Roe’s Flux Difference Splitting scheme. Integration of the equations was implicit in the time performed, until steady state was achieved, solving the linearised system of equations by the multigrid technique. The computed surface heat flux distributions were input into the thermal analysis to rebuild the thermal history of the ceramic models under the different test conditions. The thermal model is based on the solution of the unsteady energy equation in the solid, with the surface heat flux updated at each iteration to account for the energy re-emitted radiatively and for the changes in convective heat flux due to changes in surface temperature.27 The computations were carried out by using the commercial FLUENT package28 with ad hoc userdefined-functions developed in house to take into account the change of the surface temperature during the test and the radiation emitted from the surface at relatively high temperature. The thermophysical properties of the UHTC material used in this work were previously measured.3  3.3. Characterization of the UHTC models  UHTC models were cleaned with acetone in an ultrasonic bath, rinsed with distilled water, and dried in an oven at 373 K  \\x0c', '2316  Table 1  F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  Main parameters of the experimental tests: arc power (AP), average specific total  enthalpy (H0 ), maximum stagnation-point pressure (Pmax ), and duration (t). The mass flow rate was 1 g/s.  Step 1  Step 2  Step 3  Step 4  Step 5  AP (kW)  H0 (MJ/kg) Pmax (kPa) t (s)  15.9  4.5  6.8  80  18.9  5.5  7.4  60  22.4  7.3  8.1  60  25.6  8.6  8.8  60  29  10.3  9.5  230  overnight. The mass of the models was measured before and testing (accuracy 5 × 10 −2 mg). Post-test after samples were prepared for microscopy by sectioning the UHTC models into two symmetrical coupons and then polished down to a 0.25 \\u242em finish using diamond abrasives. Microstructure and composition analyses were performed by scanning electron microscopy (SEM, model S360 Leica Cambridge, UK) along with energy dispersive X-ray microanalyzer (EDX, model INCA Enegy300, Oxford Instruments, UK) on the as-exposed surfaces as well as on the polished cross-sections. A thin carbon coating was cast on the polished piece to prevent electrostatic charging of the resin (used to encase samples for polishing) during SEM observations.  4. Experimental results and discussion  4.1. Measured and derived test conditions  (Pmax )  The experiments reported here were performed with a gas flow rate of 1 g/s, and a static chamber pressure of about 200 Pa. Average specific total enthalpy (H0 ) and maximum stagnationpoint pressure ranged between 4.5-10.3 MJ/kg and 6.8-9.5 kPa, respectively. Peak cold-wall heat fluxes of 4.7 or 11 MW/m2 for the hemisphere and the cone, respectively, at the maximum H0 were calculated. Some operating conditions and derived boundary-layer-edge conditions are listed in Table 1. The ceramic models, supported on a tubular alumina holder, were moved into the hot gas flow at H0 below 5 MJ/kg. H0 was thus increased to the peak value of about 10.3 MJ/kg: ramp-up and final hold had a total duration of about 10 min. The experimental transient temperature profiles are given in Fig. 1 for both of the tested ceramic models. A transient thermal analysis (executed for the most severe step 5 in both models) compared the steady state temperature values with those obtained experimentally (Texp ). Table 2 reports the numerical averaged temperature in a 3 mm diameter spot (Tspot ) in the same area targeted by the pyrometer. In the case of the conical model, the regions facing the hot stream just in proximity of the tip have been calculated to  Table 2  Measured temperature (Texp ) and numerically computed average temperature in a 3 mm diameter spot in the same area targeted by the pyrometer  (Tspot );  an emissivity of 0.6 was input  for  the numerical simulations. The conditions  correspond to step 5 in Table 1 (steady state).  Model  Hemisphere  Cone  Texp (K)  2053  2083  Tspot (K)  2031  2089  Fig. 1. Surface temperature vs.  time (t) of  the UHTC models;  the steps  for  increasing H0 are numbered (see Table 1).  reach a surface temperature of more than 2300 K (0.6 emissivity, “un-oxidized” surface). The numerical outputs Tspot are close to the experimental values considering an emissivity of 0.6. As mentioned in Section 3.1, the pyrometer in both single and two-colour modes was used to evaluate temperature and spectral emissivity. In particular, in single colour mode the pyrometer allowed to determine the brightness of the sample at 1 \\u242em wavelength (λ) and therefore the equivalent temperature (Tb ) of a blackbody with the same spectral radiant energy at λ = 1 \\u242em. Once the real surface temperature (Texp ) was evaluated with the two-colour mode, the coefficient of spectral emissivity ελ was obtained using the following equation:  (cid:6)  (cid:4) (cid:5)  (cid:3) c2 λ  ln ελ =  1  Texp  − 1  Tb  (6)  where c2 is the Planck radiation constant. Spectral emissivity (ελ ) values at different temperatures are shown in Table 3. The reported ελ data of the UHTC specimen appear to change as its surface progressively interacts during exposure to the hot stream, changing the chemical nature of the surface under oxidation. In comparison to experiments performed in high enthalpy subsonic flow conditions (ελ ,1 \\u242em = 0.93 ), spectral emissivity was lower at the highest temperature (Table 3). This can be explained by the different dominant chemical composition of the reaction products formed at surface during oxidation. Cross-sectioned  Table 3 Measured spectral emissivity at 1 \\u242em (ελ , 1 \\u242em ) at different Temperature values were selected increasing the arc-power, according to Table 1.  temperature (T).  T (K)  1600  1700  1780  1865  2025  ελ , 1 \\u242em  Cone  0.8  0.78  0.77  0.75  0.63  Hemisphere  0.78  0.77  0.75  0.74  0.66  \\x0c', 'F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  2317  hemispherical specimens examined after exposure to high enthalpy arc-jet subsonic flow revealed the presence of a more compact external silica layer3 in comparison to the present supersonic arc-jet experiments. However, the cross-sectioned polished specimens of the present study showed a prevailing concentration of zirconia composing the external oxide layer which is responsible of a lower emissivity. These results show that the flow behavior may influence surface oxidation. One direct consequence of a reduced spectral emissivity of the surface is that the specimen reaches a stagnation-point temperature comparable to that achieved in a subsonic flow3 even though H0 is half of that applied in high pressure tests (i.e. 20 MJ/kg). The knowledge of these parameters is essential to correlate the extent of the oxidation by changing testing conditions. In the present case, the use of two surface geometries in the same flowfield has led to dissimilar boundary conditions upon the surfaces impinged by the hot stream. Whereas stagnation-point surface temperature and gas pressure were determined directly, the gas composition at the sample surface was computed through the numerical model based on the solution of the Favre-averaged Navier-Stokes equations for a mixture of reacting species in chemical and vibrational nonequilibrium. A more detailed description is reported in the next sections.  4.2. Formation and morphology of the oxidized layers  Both the UHTC specimens survived the arc-jet exposure without any optical evidence of mechanical failure (Fig. 2). Mass changes turned into net losses of 0.5 and 0.6% for the blunt hemisphere and the sharp cone, respectively. The post-test inspection at a macroscale level revealed a rather smooth contour of the oxidized surface and a whitish coloration. In the case of the cone this whitish colour was concentrated at the tip. At higher magnifications, the oxidized surfaces appear more inhomogeneous, and mostly covered by a glassy coating. The major contribution to the varying surface roughness arises from the presence of craters, which are presently interpreted as the signature of the evolution of gases during exposure. The craters were differentiated in size and shape depending on their position with respect to the stagnation point (Fig. 3). The distribution of the shear stresses (Fig. 4), computed at steady state for the same conditions reported in Table 2 (step 5) cor Fig. 2. Visual appearance of the ceramic models before (a and b) and after testing  (c and d); R: radius of curvature.  relates well with the evolution/bursting of bubbles. In fact, a localized increase in the shear stress favors an anticipated explosion of gas bubbles (see small sized craters close to the tip in Fig. 3), whilst in the regions less perturbated by the shear stresses the bubbles may evolve and grow in size in a relatively longer time (before bursting). Moving towards the colder regions of the ceramic models, the simultaneous reduction of the surface temperature (Fig. 5) and shear stresses lead the bubbles to form and eventually burst much less frequently. Looking into the cross-sectional surfaces by SEM-EDX, an interesting variety of oxidized sub-layers extending below the outermost glassy layer was revealed. The morphologies of the oxidized layers differ significantly, depending not only on the initial model’s geometry, but also on the position with respect to the stagnation point. The SEM images presented in Fig. 6, though necessarily selective, are representative of the thicknesses of the oxide layers formed. Based on the SEM-EDX observations of the layered oxide structures, an appreciable temperature gradient that developed on the exposed surfaces (Fig. 5) drove the formation of different features.  Fig. 3. Exposed surface of the sharp cone: details by SEM.  \\x0c', '2318  F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  Fig. 4. Shear stress vs. curvilinear normalized coordinate (x/L).  4.2.1. Blunt hemisphere  The entire thickness of the oxidized layer varies from 150 to about 60 \\u242em. The oxide sub-layer underlying the outermost glassy scale is characterized by a duplex morphology of the zirconia crystals: a prevailing columnar shape close to the unoxidized material, and a more globular shape nearer to the outermost surface glassy layer (Fig. 7). Also, the tiny zirconia particles which decorate the external glassy layer grow preferentially with a columnar shape. The lack of glass between adjacent columnar zirconia crystals weakens the compactness of this oxide sub-layer. On the contrary, an enrichment of glass however cements the zirconia crystals composing the upper portions of this oxide sub-layer. No obvious SiC-depleted diboride region was found, although the model’s surface rose to a temperature very close to 2100 K. Looking at the interface between the bottom of the columnar-type zirconia sub-scale and the underlying unoxidized bulk in finer details, widespread microposity is present. This porosity is believed to have nucleated during exposure to the hot plasma, when an incipient inner depletion  Fig. 5. Distribution of the surface temperature along the curvilinear normalized  coordinate (x/L). An emissivity of 0.6 was set for the calculations.  of SiC very likely began to take place due to its active oxidation. Raising the temperature, oxidation proceeded so fast that the former SiC-depleted diboride matrix readily oxidized into columnar zirconia particles, whose peculiar shape suggests the outward escape of gas species and/or transportation to the surface layer by liquid convection.13 The inner oxidation of ZrB2 was accelerated by the lack of glass which did not effectively slow the inward diffusion of oxygen. This line of reasoning is under a continuing investigation (not included herein). The presence of large zirconia crystals directly facing the external ambient (Fig. 7) should also be noted. The creation of these special features is believed to have followed the bursting of bubble(s), leaving behind zirconia crystals without any protective silica glass. A locally reduced ability to dissipate heat due to a lower emissivity and (very likely) a higher catalycity in consequence of the lack of silica glass, likely gave rise to “hot-spots” where the surface temperature jumped up at sufficient levels to rapidly sinter the underlying zirconia crystallites. Localized EDX analyses confirmed the presence of only Zr and O in such exposed enlarged crystallites.  4.2.2.  Sharp cone  On the conical sample, a glassy layer covers the major part of the external surface but appears less continuous than in the hemispherical one. The entire thickness of the reaction oxide layer varies from 190 to 50 \\u242em (Fig. 8). In contrast to the former case of the blunt hemisphere, a SiC-depleted diboride sub-region was evident, its thickness decreasing from 70 \\u242em (at the cone tip) to a complete disappearance (0.7 < x/L < 0.8, x/L the normalized curvilinear coordinate of the sample). This finding corroborates the supposition that the rate of the rise in temperature is the primary boundary condition that allows (or prevents) the creation of the SiC-depleted ZrB2 layer. The oxide sub-layer above the SiC-depleted diboride region is very close to the cone tip and has a morphology very similar to that described earlier for the blunt hemisphere, i.e. formation of both columnar and more rounded zirconia crystals. In the remaining portions of this oxide  \\x0c', 'F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  2319  Fig. 6. Cross-sectioned and polished ceramic models (SEM micrographs at same magnification).  sub-layer (far from the apex), the typical morphology consists of globular zirconia crystals cemented within a silica-based glass. The small zirconia particles which decorate the outermost glassy layer have crystallized preferentially with a regular round shape. Even if the density of the craters (i.e., the burst bubbles) appears more obvious in the present case (Fig. 9), the growth of large zirconia blocks directly facing the external ambient was not confirmed. The reason for this different behavior is not yet fully understood. A reasonable explanation may be that the more compact zirconia/glass mixed sub-scale better protected the inner bulk that still retains unoxidized diboride matrix in the SiC-depleted region. The in-depth active oxidation of SiC yields more silica, compared to the previous case. The more copious generation of silica assisted the preservation of a continuous network of glass which favors more oxidation protection and a mitigated exposure to the severe aero-thermal environment of the  zirconia skeleton composing such oxide sub-layer. It follows that conditions of quasi-steady inward diffusion of the oxidants (i.e., O and O2 ), and the outward transport of SiO and CO, are controlled by the interconnected network of silica glass. Additional investigations are in progress to determine the composition of the silica-based glass that lies upon the external face and mostly fills the gaps between the zirconia crystals constituting the oxide sub-layer. The thin, almost continuous silica glass layer that persists over the exposed surface seems to contradict some thermodynamic assessments that set conditions for volatilization29 On the other hand, an accurate knowledge of testing conditions in  Fig. 7. Cross-sectioned and polished blunt hemisphere (tip area, back-scattered  Fig. 8. Cross-sectioned and polished sharp cone (SEM micrograph): mixed zir electron SEM micrograph): outermost glassy layer decorated with tiny zirconia  conia/glass sub-layer (2) underlying the zirconia decorated external glassy layer  particles (1), oxide sub-layer (2). Large zirconia crystals are indicated.  (1), SiC-depleted diboride region (3) above the un-oxidized bulk (4).  \\x0c', '2320  F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  Fig. 9. Cross-sectioned and polished sharp cone (SEM micrograph): globular  zirconia embedded in a glass sub-layer (2) underlying the zirconia-decorated out ermost glassy layer (1), SiC-depleted diboride region (3) above the un-oxidized  bulk (4).  terms of temperature and partial pressure of the main gas species composing the reactive mixture (N2 , O, O2 N, NO) is a complex task. Experimentally, the finding of such a silica glassy layer might suggest a wider stability domain of silica under dissociated oxygen and reduced pressure. Even though the duration of the tests presented herein are representative for a re-entry mission, extended exposures seem necessary to look into the effects of time over the stability domain of silica in such a special environment. In any case, the sudden drop of the surface temperature after switching-off the aero-heating makes the very fast cooling stage not effective at all to modify chemistry and microstructure of the oxidized surface. It follows that the resulting post-test microstructures, discussed in detail above, can be considered a realistic picture of the surfaces interacting with the surrounding environment. As for the varying thickness of the SiC-depleted diboride layer, it is very likely influenced by the decreasing temperature profile along the exposed surface. Temperature aside, the oxygen partial pressure (p) is known to drive the active-to-passive oxidation transition of SiC. In the present case, the calculated patterns for pO and pO2 that vary along the exposed surface do not seem to correlate with the occurrence and extent of the SiCdepleted sub-layer. More reasonably, the expected decrease of the oxygen partial pressure through the reaction layers plays a more important role in controlling the active-to-passive transition of the SiC oxidation, and therefore the transport of diffusing species in the inward and outward direction.  4.3.  Surface properties  CFD simulations of fully catalytic versus non-catalytic heating of the test specimens suggest that the reacting surfaces must possess very low catalytic efficiency to explain the measured surface temperature transients. Surface characteristics, like efficiency for surface recombination (i.e catalycity) or emissivity, depend on both the chemical nature and the microstructure of the surface. Different from our previous experiments at atmospheric pressure and subsonic flow regime,3 the spectral emissivity  Fig. 10. Distribution of the surface temperature along the curvilinear normalized coordinate (x/L). The models have an external oxide layer 100 \\u242em thick, 1 W/mK  of thermal conductivity; a global emissivity of 0.6 was set for the calculations.     derived by simultaneous twoand one-colour radiometry measurements is in the range of 0.6-0.65. No data are available to date in order to compare the consistency of these results. Zhang and co-workers conducted tests in ground simulated reentry conditions (supersonic flow) and a surface temperature C was reached.30 No mention was made of the of 1440-1450 spectral emissivity. The authors stressed an excellent thermaloxidative and configurational stability, even though the scanned interval did not cover the ambitious range of values expected in an extreme hypersonic environment. In order to gain a more comprehensive understanding of the effects of the surface chemistry on the aero-heating, CFD simulations were built by modelling UHTC samples covered with a continuous oxide scale 100 \\u242em thick (1 W/mK of thermal conductivity, typical of compounds like zirconia and silica). Plots of the distribution of the surface temperature based on the CFD simulations (Fig. 10) show a rapid increase in temperature in the near region of the sharp apex. Such evaluations experimentally are limited by the minimum finite size of the focused spot of the radiometry determinations. As expected, the sharp cone experiences a greater rise in temperature, compared to the blunt hemisphere. Interestingly, the amplified difference in the stagnation peak temperature compared to the “un-oxidized” model (Fig. 5) agrees with the evident visual dissimilarities. Persistency and homogeneity of the external glass over the surfaces of the exposed sample influence not only the catalytic surface recombination of dissociated species (O, N) but also the solubility, diffusion and in-depth reactions of oxidants, and therefore oxidation. An update of our thermal model is in progress to take into due account such variables in order to better describe the oxidation behavior of this class of materials.  5. Conclusions  ZrB2 -SiC composites are candidates for leading-edge and control surface applications on future hypersonic aerospace vehicles, for which operating temperatures approaching 2300 K,  \\x0c', 'F. Monteverde et al. / Journal of the European Ceramic Society 30 (2010) 2313-2321  2321  or even higher, in the presence of partially dissociated oxygen are expected. This study tested samples with sharp and blunt geometries in simulated hypersonic re-entry conditions using an arc-jet plasma wind tunnel. Under the conditions of this study, measured surface temperatures near 2100 K resulted in (calculated) cold-wall heat fluxes up to 11 MW/m2 . Changes in specimen dimensions and mass were minimal. The exposures resulted in oxidation of the specimen’s surfaces, which produced an outer layer of silica-based glass and an underlying layer of mixed ZrO2 /SiO2 . The boundary conditions resulted in the development of both columnar and more rounded zirconia crystals. The measured spectral emissivity (at 1 \\u242em) had values in the range of 0.6-0.65 for temperatures above 2000 K. The persistency of a residual external silica glassy layer suggested an extension of the stability domain in such harsh conditions of temperature and pressure when oxygen is dissociated. The occurrence of graded temperature contours along the model surface was recognized as the key parameter controlling formation, extent and chemistry of the resulting oxide layered structures. The obvious formation of SiC-depleted diboride-based regions took place only in the case of the sharp cone, where the measured stagnation-point temperatures reached values of 2100 K. CFD simulations, which assumed an un-oxidized status of the exposed surface, estimated peak values close to 2300 K (cone) and 2100 K (hemisphere). Future studies are planned to update a thermal model which takes into account effects of the changing microstructure (i.e., growth of a thermally insulating scale) on the optical characteristics of this class of materials.  Acknowledgement  The authors thank Mr. D. Dalle Fabbriche (ISTEC-CNR) for technical assistance in the hot-pressing.  References  8. Monteverde F. 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Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during the     C. J Eur  oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27:2495-501.  15. Rezaie A, Fahrenholtz WG, Hilmas GE. Oxidation of zirconium diboride silicon carbide at 1500 C at a low partial pressure of oxygen. J Am Ceram Soc 2006;89(10):3240-5.     16. Zhang SC, Hilmas GE, Fahrenholtz WG.  Improved oxidation resistance  of zirconium diboride by tungsten carbide additions. J Am Ceram Soc 2008;91(11):3530-5.  17. Talmy IG, Zaykoski JA, Opeka MM. High-temperature chemistry and oxi dation of ZrB2 ceramics containing SiC, Si3N4 , Ta5 Si3 , and TaSi2 . J Am Ceram Soc 2008;91(7):2250-7.  18. Guo W-M, Zhou X-J, Zhang G-J, Kan Y-M, Li Y-G, Wang P-L. Effect  of Si and Zr additions on oxidation resistance of hot-presses ZrB2 -SiC composites with polycarbosilane as a precursor at 1500 2009;471:153-6.  C. J Alloys Compd     19. Fletcher, G. and Playez, M. Characterization of supersonic and subsonic  plasma flows, AIAA-2006-3294, San Francisco, June 2006.  20. Grinstead,  J., Driver, D.  and Raiche, G. Radial profiles of arcjet flow  properties measured with laser-induced fluorescence of atomic nitrogen,  AIAA-2003-0400, January 2003.  21. Park C, Jaffe RL, Partridge H. Chemical-kinetic parameters of hyperbolic earth entry. J Thermophys Heat Transfer 2001;15(1):76-90.  22. Hirschfelder  JO, Curtiss CF, Bird RB. Molecular  theory of gases and  liquids. New York: Wiley; 1954. p. 75-106.  23. Park C. Review of  chemical kinetics problems  for  future NASA mis sion. Part 1:  earth entries.  J of Thermophys Heat Transfer 1993;7(3):  1. Marschall J, Pejakovic DA, Fahrenholtz WG, Hilmas GE, Zhu S, Ridge J,  385-98.  et al. Oxidation of ZrB2 -SiC ultrahigh-temperature ceramic composites in dissociated air. J Thermophys Heat Transfer 2009;23(2):267-78.  24. Millikan RC, White DR. Systematics of vibrational Phys 1963;39(12):3209-13.  relaxation. J Chem  2. Zhang X, Weng L, Han J, Meng S, Han W, Preparation.  thermal ablation  25. Park C. Nonequilibrium hypersonic aerothermodynamics. John Wiley &  behavior of HfB2 -SiC based ultra-high-temperature ceramics under severe heat conditions. Int J Appl Ceram Technol 2009;6(2):134-44.  3. Monteverde F, Savino R. Stability of ultra-high-temperature ZrB2 -SiC ceramics under simulated atmospheric re-entry conditions. J Eur Ceram Soc 2007;27:4797-805.  4. Monteverde F. Ultra-high temperature HfB2 -SiC ceramics consolidated by hot-pressing and spark plasma sintering. J Alloys Compd 2007;428/12:197-205.  5. Monteverde F. Hot pressing of hafnium diboride aided by different sinter additives. J Mater Sci 2008;43:1002-7.  Sons; 1990.  26. Suslov, O. N. and Tirskiy, G. A. Kinetics of the recombination of nitro gen atoms on high temperature reusable surface insulation in hypersonic  thermo-chemical non-equilibrium flow; 2004. In: Hunt, J. J., editor, Pro ceedings of  the 2nd European Symposium on Aerothermodynamics  for  Space Vehicles, ESTEC, Noordwijk, The Netherlands, 21-25 November  1994, European Space Agency (ESA), Paris, 1995, p. 413-9.  27. Savino R, De Stefano Fumo M, Paterna D, Serpico M. Aerothermodynamic  study of UHTC-based thermal protection systems. Aerospace Sci Technol 2005;9(2):151-60.  6. Sciti D, Monteverde F, Guicciardi S, Pezzotti G, Bellosi A. Microstructure  28. Fluent 6.2 User’s Guide. Lebanon, NH, USA: Fluent Inc; 2004.  and mechanical properties of ZrB2 -MoSi2 ceramic composites produced by different sintering techniques. Mater Sci Eng A 2006;434:303-9.  7. Tang S, Deng J, Wang S, Liu W. Comparison of  thermal and ablation  behavior of C/SiC composites 2009;51:54-61.  and C/ZrB2 -SiC composites. Corr Sci  29. Fahrenholtz WG. Thermodynamic analysis of ZrB2 -SiC oxidation: formation of a SiC-depleted region. J Am Ceram Soc 2007;90(1):143-8.  30. Zhang X, Hu P, Han J, Meng S. 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},{
  "_id": 209,
  "PDF": "Preparation and characterization of high-temperature resistant ZrC-ZrB2 nanocomposite ceramics derived from single-source precursor.pdf",
  "Text": "[\"Materials and Design 117 (2017) 257-264  Contents lists available at ScienceDirect  Materials and Design  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t d e s  Preparation and characterization of high-temperature resistant ZrC-ZrB2 nanocomposite ceramics derived from single-source precursor  Shugang Chen, Yanzi Gou ⁎, Hao Wang, Ke Jian, Jun Wang  Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, People's Republic of China  H I G H L I G H T S  G R A P H I C A L  A B S T R A C T  • Precursors were synthesized by reactions of allyl amine, ZrCl4, borane and the Grignard reagent (allyl-MgCl). • Ceramization processes of the ZrNCB precursors were analyzed and ZrC-ZrB2 ceramics were obtained. • The obtained nanocomposite ceramics have good high-temperature stability up to 2000 °C with dense and nano scaled structure. • The ZrC-ZrB2 ceramics also have excellent stability at 1700 °C under air due to the oxidation layer formed on the surface.  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 28 July 2016 Received in revised form 8 December 2016 Accepted 15 December 2016 Available online 16 December 2016  Keywords:  Precursor  Topic:  ZrC-ZrB2 Ceramization High-temperature resistance Oxidation resistance  1. Introduction  In order to prepare ZrC-ZrB2 nanocomposite ceramics with high temperature stability from single-source precursors, the reactions of ZrCl4, allyl amine, allyl-MgCl and borane were employed to synthesize suitable precursors. The meliorated ZrNCB precursor had favorable solubility and higher ceramic yield (68.8 wt%), which was then transformed to ZrC-ZrB2 ceramics at different temperatures. The main phase of the ceramics was amorphous at 1000 °C. As the temperature increased, the amorphous phases started to crystallize, resulting in the appearance of ZrC and ZrB2 peaks. The content of ZrC and ZrB2 phases was relatively high in the nanocomposite ceramics at 1600 °C (over 80 wt%). The ZrC-ZrB2 nanocomposite ceramics showed good high temperature stability up to 2000 °C. Oxidation test indicated that an oxide layer was formed on the ceramic surface at 1700 °C under air to prevent further oxidation of the ceramics. With excellent high temperature and oxidation resistances, the ZrC-ZrB2 nanocomposite ceramics have potential application for advanced rocket engines and nuclear industry. © 2016 Elsevier Ltd. All rights reserved.  Due to excellent high temperature stability and oxidation resistance, Zr-based ceramics, such as ZrC-ZrB2, ZrC-SiC and ZrB2-SiC [1-3], have  ⁎  Corresponding author. E-mail addresses: y.gou2012@hotmail.com (Y. Gou), wangjun_cfc@nudt.edu.cn (J. Wang).  http://dx.doi.org/10.1016/j.matdes.2016.12.041 0264-1275/© 2016 Elsevier Ltd. All rights reserved.  potential application in the fields of advanced rocket engines and nuclear industry [4-6]. It has been reported that the properties could be greatly improved by reducing the grain size of the Zr-based ceramics to the nano scale [7-10]. However, the preparation of nanocomposite ceramics using common methods has been frustrated by difficulties to get dense structures and to control grain growth at high temperatures [11,12]. Recently, the precursor derived method has been proved to be a favorable tool to fabricate nanocomposite ceramics with complicated shapes [13-  \\x0c\", '16], as single-source precursor can be modulated to achieve unique phase composition, microstructure and property of the derived ceramics [17,18]. In order to synthesize suitable precursors for Zr-based nanocomposite ceramics, several strategies have been developed, including: (a) Formation of Zr  \\x0c', '(TG-MS) through employing a thermal analysis device (TGA/DSC 2 Mettler Toledo, Switzerland) coupled with a mass spectrometer (QMG700, Inficon, Germany). X-ray diffraction (XRD) studies were carried out with a Bruker AXS D8 Advance diffractometer (Bruker, Germany) with Cu Kα radiation (k = 1.54178 Å). The morphology of the samples was examined with a scanning electron microscope (SEM; HITACHI S-4800, Japan). The crystalline behavior was observed using a transmission electron microscope (TEM; Titan G2 60-300, Holland). The oxygen and nitrogen contents, as well as the content of carbon of the derived ceramics were determined by using a Horiba oxygen/nitrogen analyzer EMIA-820 (Horiba, Japan) and a Horiba carbon/sulfur analyzer EMIA-320 V. The zirconium content was determined using a VistaMPX inductively coupled plasma emission spectrometer (Variant, America). The contents of boron in the samples were determined by utilizing the neutralization titration method.  was maintained at the pyrolytic temperature for 2 h before finally being cooled down to room temperature. Pyrolytic temperatures of 1000, 1200, 1400 and 1600 °C were chosen to investigate the ceramization of the precursor and crystallization process of the derived ceramics. The ceramics derived from ZrNCB-2 precursor at 1400 °C was heated to 1800 and 2000 °C, respectively, and held for 1 h to measure the high temperature resistance.  2.4. Oxidation test of the precursor  3. Results and discussion  3.1. Synthesis and characterization of the ZrNCB precursors  As illustrated in Scheme 1, firstly, Zr  To investigate the oxidation resistance of the ceramics, the ceramics derived from ZrNCB-2 precursor at 1400 °C were heated to 1200 °C at a rate of 10 °C min−1 under air by thermogravimetric-differential scanning calorimeter (TG-DSC). Moreover, in order to further evaluate the oxidation resistance at higher temperature, a ceramic specimen was prepared from the ZrNCB-2 precursor. The precursor was treated at 300 °C for 1 h and then formed into wafers with a diameter of 10 mm. The wafers were sintered at 1600 °C for 1 h under argon atmosphere. These ceramic wafers were then heated to 1700 °C at a rate of 5 °C min−1 and kept at the temperature for 30 min under air in a muffle furnace to test the oxidation resistance of the precursor derived ceramics at high temperature.  2.5. Characterization  FT-IR spectra were recorded with a Nicolet Avator 360 apparatus (Nicolet, America) using KBr pellets. 1H NMR data were collected in DMSO-d6 with a Bruker AV-400 spectrometer (Bruker, Germany) operating at 75.46 MHz. The polymer-to-ceramic transformation of the precursor was studied by thermogravimetric analysis-mass spectroscopy  \\x0c', '3.2. Ceramization of the ZrNCB precursors  TG-MS measurement was used to explore the organic-to-inorganic process of the precursors, as shown in Fig. 3. The ceramic yield of the ZrNCB-1 precursor at 1000 °C is about 46.0%. The release of alkanes and NH3 is the main reason for the low ceramic yield of this precursor. The weight loss around 250 °C is due to the release of CH3CH3 through the reaction of B  \\x0c', \"S. Chen et al. / Materials and Design 117 (2017) 257-264  261  Table 3 Chemical composition of the ceramics pyrolyzed at 1600 °C from bulk chemical analysis.  Composition (wt%)  Sample  ZrC-ZrB2-1 ZrC-ZrB2-2  Zr  55.32 58.25  C  34.10 30.05  B  3.85 5.83  N  6.57 5.01  O  0.16 0.08  Phases (wt%) Zr(C, N):ZrB2  32.38: 40.39 30.23: 52.05  Å 0CZrNCB→ZrðC; NÞ þ ZrB2 þ BN þ C 1200e160  ð8Þ  3.3. High temperature resistance of the derived ZrC-ZrB2 ceramics  In order to investigate the high temperature resistance, the ceramics derived from ZrNCB-2 precursor at 1400 °C were heated to 1800 and 2000 °C, respectively. Fig. 6(a) shows the XRD curves of the ceramics treated at high temperatures. When temperature increases to 1800 and 2000 °C, a slight shift (0.1-0.3°) of ZrC peak to a lower diffraction angle could be noticed due to the transformation of Zr(C, N) to ZrC and ZrB2, as stated in Eqs. (9) and (10). The composition of Zr(C, N) ceramics could be estimated from the position of the reflections using the Vegard's rule, which is shown in the supporting information. According to the XRD patterns, the average grain size of ZrC in the ceramics is 32.6 nm at 1800 °C, which is calculated from the ZrC (200) half-peak width by the Scherrer formula. According to the literature, the grain  size of ZrC increases rapidly above 1400 °C and can reach up to several micrometers over 1600 °C [31]. Therefore, it is indicated that the presence of ZrB2 could prevent the unfavorable growth of crystalline ZrC.  ZrðC; NÞ þ C→ZrC þ N2  ZrðC; NÞ þ BN→ZrB2 þ N2 þ C  ð9Þ  ð10Þ  The morphologies of the ZrC-ZrB2 ceramics at 1400, 1800 and 2000 °C are illustrated by Fig. 6. As seen from Fig. 6(b), (c) and (d), with increasing temperature, the ZrC-ZrB2 ceramics become densified on a macroscale. This dense structure is expected to be beneficial to the mechanical properties for application as the matrix of CMCs under high temperature of 2000 °C, which should be better than the ceramics obtained from sol-gel method [32]. The microstructure of the ZrC-ZrB2 ceramics derived at 1800 °C is further revealed by TEM and HRTEM in Fig. 7. It is noticeable that the ceramics have a uniform distribution with nano scaled crystallites. The elemental mappings also demonstrate the homogeneous dispersion of phases and elements. The nanocomposite structure of the derived ceramics is essential for the densified appearance [33]. Table 4 shows mass changes and compositions of the ZrC-ZrB2 ceramics after treatment at 1800 and 2000 °C. The ceramics undergo little mass changes due to the transformation of Zr(C, N) to ZrC at high temperatures (− 1.23% and − 1.25% for 1800 °C and 2000 °C, respectively). Therefore, it is apparently proved that the ZrC-ZrB2 nanocomposite ceramics have excellent stability at high temperature up to 2000 °C. In the ceramics the amount of high temperature phases (ZrC and ZrB2) is  Fig. 6. XRD curves (a) and SEM images of ZrC-ZrB2 ceramics at 1400 (b), 1800 (c) and 2000 °C (d).  \\x0c\", '262  S. Chen et al. / Materials and Design 117 (2017) 257-264  Fig. 7. TEM (a), HRTEM (b), HADDF scanning image (c) and elemental mappings (d) of ZrC-ZrB2 ceramics obtained at 1800 °C.  higher than 80 wt%, which could guarantee the excellent properties of the ceramics as ultra-high temperature resistant materials.  3.4. Oxidation resistance of the derived ZrC-ZrB2 ceramics  The ceramics derived from ZrNCB-2 precursors at 1600 °C were tested by TG-DSC under air at low temperatures (20-1200 °C). It can be seen from Fig. 8(a) that the sample weight remains unchanged with only a slight peak of 1.89 wt% weight gain at about 650 °C, which can  Table 4 Mass changes and chemical composition of the ceramics treated at 1800 °C and 2000 °C.  Sample  Mass change (%)  1800 °C − 1.23 2000 °C − 1.25  Composition (wt%)  Zr  C  B  N  O  Phases (wt%) Zr(C, N):ZrB2:carbon  66.82 70.95  20.94 14.79  9.72 10.51  2.47 3.18  0.05 0.09  30.16:52.33:17.51 31.28:57.12:11.60  be attributed to the oxidation of ZrC. Obvious exothermic peaks at 589.6 °C and 664.6 °C can be observed in the DSC curves, corresponding to the oxidation of ZrC and residual carbon. When temperature gets higher than 1000 °C, the sample weight starts to increase slowly (4.78 wt% weight gain up to 1200 °C) due to the oxidation of ZrB2 [34]. The ceramics were also tested in a muffle furnace at a higher temperature of 1700 °C. The weight and the diameter gains of the wafer specimen are 2.01 wt% and 0.02%, respectively, which could be attributed to the evaporation of B2O3 at high temperature [35]. The low weight and diameter gains illustrate excellent oxidation resistance of the ceramics. It can be noticed from the SEM image in Fig. 8(b) that the oxide layer of the ceramics could be divided to two parts. The outer part is a thin layer of about 3 μm formed on the surface of the specimen. XRD and EDS results (in supporting information) show that it is mainly composed of B2O3. The inner part is a mixed oxidized phases of B2O3, ZrO2, ZrC and ZrB2. It can be seen that this layer is compact, which would act as the oxidation barrier to protect the ceramics from further oxidation.  \\x0c', 'S. Chen et al. / Materials and Design 117 (2017) 257-264  263  Fig. 8. (a) TG-DSC curves in air up to 1200 °C and (b) SEM image of ZrC-ZrB2 ceramics treated at 1700 °C in a muffle furnace.  4. Conclusions  Soluble precursors for ZrC-ZrB2 nanocomposite ceramics were prepared by using ZrCl4, allyl amine, borane and allyl-MgCl. The TG-MS measurements proved that the usage of Grignard reagent was favorable for higher ceramics yield (68.8 wt%). The derived ceramics at 1600 °C were composed of nano scaled ZrC and ZrB2 crystallites with a high content (30.23 wt% for ZrC and 52.05 wt% for ZrB2, respectively). Annealing under argon atmosphere confirmed that the derived ZrC-ZrB2 ceramics had excellent high temperature resistance up to 2000 °C. Oxidation tests showed that the nanocomposite ceramics were stable at high temperatures due to the formation of oxide layer (B2O3 and ZrO2) on the surface of the ceramics. The excellent high temperature and oxidation resistances make ZrC-ZrB2 nanocomposite ceramics appealing candidates for application in ceramic matrix composites.  Acknowledgments  The authors thank the National Natural Science Foundation of China (grant no. 51302313), the financial support from Postdoctoral Science Foundation of China (2014M552685), Aid program for Innovative Group of National University of Defense Technology and the Aid Program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province.  References  [3]  [1]  S. Chen, Y. Gou, H. Wang, J. Wang, Fabrication and characterization of precursor-derived non-oxide ZrC-SiC multiphase ultrahigh temperature ceramics, J. Eur. Ceram. Soc. 36 (2016) 3843-3850. [2] H . Wang , B . Gao , X . Chen , J . Wang , S . 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Hoffman, Thermal stability of refractory carbide/boride composites, Mater. Chem. Phys. 74 (3) (2002) 272-281. [35] W.M. Guo, G.J. Zhang, Oxidation resistance and strength retention of ZrB2-SiC ceramics, J. Eur. Ceram. Soc. 30 (2010) 2387-2395.  [34]  \\x0c']"
},{
  "_id": 210,
  "PDF": "Preparation and properties of dense ZrB 2 composite reinforced by elongated SiC and Al 3 BC 3 grains.pdf",
  "Text": "['Received: 9 April 2019   DOI: 10.1111/ijac.13265    |   Revised: 9 April 2019   |   Accepted: 25 April 2019  O R I G I N A L A R T I C L E  Preparation and properties of dense ZrB2 composite reinforced  by elongated SiC and Al3BC3 grains  Zhibo\\xa0Chen1  Rui\\xa0Zhang2,3 Sea‐Hoon\\xa0Lee5  |   Xiaotong\\xa0Zhao2 |   Hongliang\\xa0Xu2   |   Hailong\\xa0Wang1,2 |   Gang\\xa0Shao2  |   Wen\\xa0Liu2  |   Bingbing\\xa0Fan2  |   Hongxia\\xa0Lu2  |   Yanhui\\xa0Chu4  |    |     Abstract  Dense ZrB2‐SiC‐Al3BC3 ultra‐high temperature ceramic composite was fabricated  by hot pressing sintering at 1900°C for 1\\xa0 hour under a pressure of 20\\xa0 MPa using  Zirconium diboride (ZrB2) as the raw material and a powder mixture of SiC, B4C,  Al, and carbon as the sintering additive. Al and B4C underwent in situ reaction with  carbon powder to produce Al3BC3, which promoted the densification of ZrB2 ceramic. SiC grains were found to be elongated during sintering. The ZrB2‐SiC‐Al3BC3  composite exhibited excellent mechanical properties, such as high flexural strength  of 589\\xa0±\\xa0 147\\xa0MPa and fracture toughness of 7.81\\xa0±\\xa0 1.09\\xa0MPa\\xa0m1/2. Oxidation behavior of the ZrB2‐SiC‐Al3BC3 composite was studied in air at 1500°C for 1\\xa0hour. A  continuous layer of oxides consisting of a mixture of SiO2, Al2SiO5, and Al2O3 was  formed on the surface of the ZrB2‐SiC‐Al3BC3 composite. This layer of oxides efficiently prevented oxygen from diffusing into the specimens during oxidation, which  improved the oxidation resistance of the ZrB2 ceramics.  K E Y W O R D S  elongated SiC, hot pressing, mechanical properties, oxidation resistance  1Henan Province Industrial Technology  Research Institute of Resources and  Materials, Zhengzhou University,  Zhengzhou, China 2School of Materials Science and  Engineering,\\xa0Zhengzhou University,  Zhengzhou, China 3Henan Key Laboratory of  Aeronautical Materials and Application  Technology,\\xa0Zhengzhou Institute of  Aeronautics, Zhengzhou, China 4School of Materials Science and  Engineering,\\xa0South China University of  Technology, Guangzhou, China 5Division of Powder and Ceramics  Research,\\xa0Korea Institute of Materials  Science, Changwon, Republic of Korea  Correspondence  Hailong Wang, Gang Shao,\\xa0School of  Materials Science and Engineering,  Zhengzhou University, Zhengzhou 450001,  China. Emails: 119whl@zzu.edu.cn;Gang_shao@ zzu.edu.cn  Funding information  National Natural Science Foundation of  China, Grant/Award Number: 51772275  1   |   I N T RO D U C T I O N  Zirconium diboride (ZrB2) has high melting temperature  (3245°C), excellent strength at elevated temperatures, and  high thermal and electrical conductivities, which make it an  ideal candidate for applications in thermal protection systems  and propulsion systems for hypersonic aerospace vehicles.1‒4  However, the application of ZrB2 ceramic is limited due to  its poor sinterability, low fracture toughness, poor oxidation   resistance, and low thermal shock resistance. The poor sinterability of ZrB2 ceramic is due to the presence of strong  covalent bonds in its structure, high melting point, and low  self‐diffusion coefficient. Therefore,  the densification of  ZrB2 ceramic material requires high temperature and external pressure. Recently, several field‐assisted sintering methods  including spark plasma sintering (SPS),5 microwave  sintering,6 and hot pressing (HP)7 have been proposed to  densify ZrB2 and ZrB2‐based ceramics.7 Moreover, the oxide   2190   |   wileyonlinelibrary.com/journal/ijac   Int J Appl Ceram Technol. 2019;16:2190-2196.  © 2019 The American Ceramic Society           \\x0c', 'impurities attached to ZrB2 particles and the large particle  size of starting powders hinder the densification of ZrB2.8‒10  Therefore, reducing the size of initial particles or adding sintering additives can promote the densification and enhanced  mechanical properties of ZrB2. Fahrenholtz et al11 stated that  the reaction of B4C or WC with ZrO2 could reduce the sintering temperature of ZrB2. Sha et al12 claimed that short carbon fibers could toughen ZrB2‐based ceramics. The f lexural  strength and fracture toughness of composite were 458\\xa0MPa  and 6.9\\xa0MPa\\xa0m1/2, respectively. Addition of proper amounts  of SiC particles can both improve the sinterability and enhance the oxidation resistance of ZrB2.7,13 Hwang et al14 reported that the addition of small size SiC particles imparted  better oxidation resistance and improved the sinterability of  ZrB2 ceramics. ZrB2 matr ix composites with textured microstructure  exhibit anisotropic mechanical proper ties. Guo15,16 claimed  that a textured ZrB2‐ZrC‐C ceramic showed high strength  and high fracture toughness. ZrB2 grains can or ient and  ar range themselves in prefer red directions to acquire a  textured ceramic under external gradient f ields, such as  mechanical f ield, electr ic f ield, magnetic f ield, and temperature f ield. Rod‐like or lath‐like grains can rotate and  ar range in an orderly two‐dimensional (2D) or 1D manner  under unidirectional pressure dur ing sinter ing or subsequent hot forging, thus resulting in a textured ceramic. Liu  et al17 obtained textured ZrB2‐MoSi2 ceramic composites  with or ientation degree of 0.59 by hot forging. The f lexural  strength of the obtained composite was 871\\xa0±\\xa0103\\xa0MPa. Ni  et al18 sintered ZrB2‐20 vol% SiC composites with textured  microstructure by SPS at 1900°C. The fracture toughness  of the obtained composite was 3.9\\xa0±\\xa00.12\\xa0MPa\\xa0m1/2. AKin  et al19 also sintered ZrB2‐50 mass% SiC textured ceramics  by SPS at 1900°C. The toughness of the obtained composite was 4.1\\xa0MPa\\xa0m1/2. Recently, we reported the use of sintering additives based  on Al-B-C system (Al mixed with B4C and carbon black  powders) to efficiently enhance the sinterability of ZrB2‐ based ceramics. Furthermore, Al-B-C system can improve  the toughness and strength of ZrB2 ceramics.20 The in situ  reaction product Al3BC3 can not only reduce the size and  amount of pores in the specimen, but also the oxide impurities due to the reaction between carbon and the oxides on  ZrB2 particles surface.7,21 Furthermore, the oxidation products formed a protective  layer which improved the oxidation resistance even in water  vapor atmosphere.22 Based on the above mentioned discussion, the main objective of this study was to prepare a dense  ZrB2 composite sintered by HP method using both Al-B-C  system and SiC as sintering additives. Moreover, the microstructure and oxidation resistance of the sintered specimens  were systematically investigated. Effects of the microstructure on the mechanical properties were also evaluated.  |   2   2.1   E X P E R I M E N TA L  |   2191  | Material preparation  The starting materials used in this study were ZrB2 powder (1‐2\\xa0 μm, purity ≥99%), Al (11\\xa0 μm, purity ≥99.7%),  B4C (<10\\xa0 μm, purity ≥99%), and carbon black powder  (60‐80\\xa0m2\\xa0g−1, purity ≥99.5%), and they were all purchased  from Alfa‐Aesar, MA, USA. The SiC particles (0.5\\xa0μm, purity ≥99.5%; Alfa‐Aesar, MA, USA) existed in the form of  β‐SiC. Al, B4C, and carbon black powders were used to synthesize Al3BC3, which formed the Al-B-C system. Powders  of 95 wt% (ZrB2‐30 vol% SiC) + 5 wt% Al3BC3 (the molar  ratio of Al, B, and C was 3:1:3) were ball milled in a polyethylene bottle for 4\\xa0hour. ZrO2 balls and ethanol were used as  the grinding media. Further, the mixtures were homogenized  again by ultrasonic dispersion for 30\\xa0minutes, and then dried  by rotary evaporation. The powder mixtures were loaded into  a graphite mold (inner diameter 30\\xa0mm) lined with graphite  foil and sintered by HP at 1900°C under a pressure of 20\\xa0MPa  for 1\\xa0hour in Ar atmosphere.  2.2   |   Characterization  The  sintered  samples  were  cut  into  test  bars  (3\\xa0 mm\\xa0 ×\\xa0 2\\xa0 mm\\xa0 ×\\xa0 25\\xa0 mm) by laser cutting. After surface  grinding and polishing, three‐point bending test was used  to measure the bending strength of sintered specimens with  a span of 10\\xa0 mm, and the crosshead speed of testing was  0.5\\xa0mm\\xa0min−1. Archimedes’ method was used to measure the  bulk density of sintered specimens. The rule of the mixture  was used to calculate the theoretical densities of the sintered  specimens. The densities of ZrB2, SiC, and Al3BC3 are 6.09,  3.20, and 2.66\\xa0g\\xa0 cm−3, respectively. Vickers’ hardness (Hv)  and the fracture toughness of sintered specimens were measured by Vickers indention method (TH700) with a 1\\xa0kg load  applied for 15\\xa0seconds, according to the Anstis formula23:  Hv = [1.854P]\\x01d2  KIC = \\x020.016 (E∕Hv)1∕2 P\\x02\\x02c3∕2  (1)  (2)  Oxidation studies were conducted in a tube furnace by  exposing the specimens to stagnant air at 1500°C for 1\\xa0hour.  The cubic specimens were placed on an alumina plate with  minimal contact area to avoid interaction at high temperature.  The heating rate was 10°C\\xa0min−1. Mass of the specimens before and after oxidation was measured using an electronic  precision balance with mg sensitivity. Microstructure of sintered specimens was characterized  by scanning electron microscopy (SEM, JSM‐6700F, Japan;  Zeiss/Auriga FIB, Germany)  and  transmission  electron   CHEN Et al.      \\x0c', '2192   |   F I G U R E 1   (A) Scanning electron  microscopy micrograph of SiC starting  powder and (B) X‐ray diffraction pattern of  the ZrB2‐SiC‐Al3BC3 composite sintered by  hot pressing at 1900°C  microscopy (TEM, JEM 2100 F, JEOL, Tokyo, Japan) along  with energy dispersive spectroscopy (EDS) for elemental  analysis. Phase constituents of the sintered specimens were  identified by X‐ray diffraction (XRD, D‐mAX‐3B, Japan)  using Cu Kα radiation.  black powder reacted with the oxide impurities (eg ZrO2)  attached to ZrB2 particles to promote the densification of  ZrB2.10 Subsequently, possible reactions between Al, B4C,  and carbon black powder occurred to form Al3BC3 at high  temperature as follows:  R E S U LT S A N D D I S C U S S I O N  9Al (l) + 2B4C (s) = 2Al3BC (s) + 3AlB2 (s) (813 C)  |   3   3.1   | Microstructural analysis  Figure 1B presents the XRD patterns of the sintered specimens, revealing that the main phase is ZrB2. Moreover, low  intensity peaks of β‐SiC and Al3BC3 are also detected. This  indicated that phase transition of β‐SiC and formation of  Al3BC3 occurred during sintering. Figure 2A presents the  surface microstructure of the ZrB2‐SiC‐Al3BC3 composite.  Figure 2B exhibits the internal structure of the ZrB2‐SiC‐ Al3BC3 composite. It was observed that the microstructure  was dense and uniform, and almost no pores existed in the  samples. Moreover, the measured relative density of the sintered ZrB2‐SiC‐Al3BC3 composite was 99.3%. These results  clearly indicated that the densification was practically complete. The densification of ZrB2 was enhanced mainly due to  the reactions between the sintering additives and the oxides  on the surface of ZrB2 particles. First, metal Al melted and  coated the other solid additives, which promoted their dissolution into the Al melt when the sintering temperature reached  above 660°C. Further, the liquid Al filled into the gaps between SiC and ZrB2 particles to promote the rearrangement  and densification. With the increase in temperature, carbon   (A)  (B)  (3)  (4)  (5)  4Al3BC (s) + 8AlB2 (s) + 16C (s)  = 5Al4C3 (s) + 5B4C (s) (1095 C)  3Al4C3 (s) + B4C (s) + 2C (s) = 4Al3BC3 (s) (1180 C)  These processes also enhanced the densification of ZrB2.  In order to observe the phase distribution in sample, elemental mapping analysis was performed (as shown in Figure 3).  Clearly, the distribution of various particles in sample was  relatively uniform. Gray grains of ZrB2 with average size  of 3‐10\\xa0 μm were observed. Moreover, some white rod‐like  grains were also observed, which were identified as SiC.  However, these SiC grains were larger than the starting SiC  particles (shown in Figure 1A). This indicated that the rod‐ like SiC grains were formed due to the growth of fine SiC  particles during the sintering process. The TEM  image shown  in Figure 4 also conf irms  the presence of rod‐like white SiC grains. Moreover, it  clearly reveals  the boundary of par ticles among specimen. Moreover, special zones exist between the ZrB2 and  SiC grains, which may be the glass phase formed by the   F I G U R E 2   (A) Typical  microstructure of specimen and (B)  scanning electron microscopy micrograph  of internal structure of the ZrB2‐SiC‐Al3BC3  composite  CHEN Et al.    \\x0c', 'F I G U R E 3   Elemental mapping  analysis of the ZrB2‐SiC‐Al3BC3 composite  sintered by hot pressing at 1900°C [Color  figure can be viewed at wileyonlinelibrary. com]  |   2193  F I G U R E 4   TEM images of different  areas of specimen and energy dispersive  spectroscopy data of specified spots [Color  figure can be viewed at wileyonlinelibrary. com]  oxidization of the star ting mater ials. However, no peaks  of oxides are detected as shown in Figure 1B. Some B4C  reacted with oxide product ZrO2 attached to ZrB2 to form  amorphous B2O3. Moreover, the O element and B element  could be detected by TEM as shown in Points B and C   of Figure 4. Thus it was speculated that the glass phase  was mainly B2O3. This could cause  the rear rangement  and growth of SiC grains. The development of elongated  SiC grains with large aspect ratio was related to the β-α  phase  transformation dur ing  liquid phase sinter ing.24,25   CHEN Et al.      \\x0c', '|   2194   (A)  (B)  F I G U R E 5   Scanning electron  microscopy images of the crack path on the  surface of the ZrB2‐SiC‐Al3BC3 composite  sintered by hot pressing at 1900°C. A, The  crack deflection and bridge. B, The crack  bridge  The β-α phase transformation may occur at temperatures  below 1900°C with Al-B-C additives.26 According to the  data presented in Figure 1B, it is evident that α‐SiC was  generated dur ing the heating process. In the early stage of  transformation, the solubility of β‐SiC in liquid phase was  higher than that of α‐SiC, which promoted grain growth.  With the progress in the transformation, the α‐SiC grains  sur rounded by β‐SiC grains were in a higher state of supersaturation, which caused elongation of the grains.27  in the sintered specimen. Noteworthy, the cracks def lected  around the elongated SiC, and caused the crack surface to  be no longer perpendicular to the applied stress. The expansion of the crack was subjected to resistance. On the other  hand, elongated SiC that existed behind the crack tip formed  an intact bridging. This bridging zone reduced the width of  crack‐opening. Therefore, more energy was consumed when  the crack surfaces were separated from each other. This result  indicated that the existence of elongated SiC grains enhanced  the toughness of sintered specimens.  3.2   | Mechanical properties  The hardness of the sintered ZrB2‐SiC‐Al3BC3 composite in  this study was 17.7\\xa0 ±\\xa0 0.8 GPa, which is lower than that of  ZrB2 (23 GPa)11 and SiC (22.9 GPa).28 The main reason is  that the hardness of Al3BC3 is only 13.9 GPa.29 The flexural  strength of the ZrB2 composite was 589\\xa0±\\xa0147\\xa0MPa, which is  higher than that of monolithic ZrB2.9 The fracture toughness of the ZrB2‐SiC‐Al3BC3 composite was 7.81\\xa0 ±\\xa0 1.09\\xa0 MPa\\xa0 m1/2. SEM images were used to  observe the extension path of Vickers‐indentation‐induced  cracks in the specimens (as shown in Figure 5). A type of  “self‐reinforced microstructure” consisting of  large elongated SiC grains and relatively fine elongated SiC grains  (black constituents) was developed from the β‐SiC starting powder.30 Crack bridging and def lection were observed   3.3   | Oxidation resistance of the composites  The increase in mass of the specimens was 2.55% after oxidation at 1500°C for 1\\xa0 hour. In oxygen‐containing environment, the composites may react with oxygen as follows31:  2SiC (s) + 3O2 (g) = 2SiO2 (g) + 2CO (g)  2ZrB2 (s) + 5O2 (g) = 2ZrO2 (s) + 2B2O3 (l)  B2O3 (l) = B2O3 (g)  2Al3BC3 + 12O2 = 3Al2O3 + B2O3 (l) + 6CO2 (g)  (4)  (5)  (6)  (7)  (A)  (B)  F I G U R E 6   (A) Scanning electron microscopy images of the cross‐section perpendicular to the oxide layer and (B) Microstructure and energy  dispersive spectroscopy data of specimen surface after oxidation [Color figure can be viewed at wileyonlinelibrary.com]  CHEN Et al.    \\x0c', '|   2195  7.81\\xa0 ±\\xa0 1.09\\xa0 MPa\\xa0 m1/2, respectively. Crack bridging and deflection caused by elongated SiC grains were the major mechanisms responsible for the toughening of the ZrB2‐SiC‐Al3BC3  composite. The oxidation resistance of ZrB2 was improved  due to the continuous layer of oxides formed on the surface of  ZrB2‐SiC‐Al3BC3 composite, and this layer was composed of  Al2O3, Al2SiO5, and SiO2. This layer efficiently prevented the  diffusion of oxygen into the specimens during oxidation.  AC K N OW L E D G E M E N T  The authors greatly acknowledge the financial support from  the National Natural Science Foundation of China (Grant No.  51772275).  O RC I D  Xiaotong Zhao\\xa0 Hailong Wang\\xa0 Rui Zhang\\xa0 Yanhui Chu\\xa0 Sea‐Hoon Lee\\xa0   https://orcid.org/0000-0003-0746-3799   https://orcid.org/0000-0003-2993-8396   https://orcid.org/0000-0003-1836-1245   https://orcid.org/0000-0001-6158-7501   https://orcid.org/0000-0002-7276-3264   R E F E R E N C E S   1. Fahrenholtz WG, Hilmas GE, Zhang SC, Zhu S. Pressureless  Sintering of Zirconium Diboride: particle size and additive effects.  J Am Ceram Soc. 2008;91:1398-404.  2. Guo WM, Tan DW, Zhang ZL, Xie H, Wu LX, Lin HT. Synthesis  of fine ZrB2 powders by new borothermal reduction of coarse ZrO2  powders. J Ceram Int. 2016;42:15087-90.  3. Wang HL, Fan BB, Feng L, Chen DL, Lu HX, Xu HL, et al. The  fabrication and mechanical properties of SiC/ZrB2 laminated ceramic composite prepared by spark plasma sintering. J Ceram Int.  2012;38:5015-22.  4. Bai LY, Ni SL, Jin HC, He JP, Ouyang YG, Yuan FL. ZrB2 powders  with low oxygen content: synthesis and characterization. Int J Appl  Ceram Technol. 2018;15:508-13.  5. Balak Z, Asl MS, Azizieh M, Kafashan H, Hayati R. Effect of different additives and open porosity on fracture toughness of ZrB2‐SiC‐ based composites prepared by SPS. J Ceram Int. 2017;43:2209-20.  6. Cao YN, Zhang HJ, Li FL, Lu LL, Zhang SW. Preparation and  characterization of ultrafine ZrB2-SiC composite powders by a  combined sol-gel and microwave boro/carbothermal reduction  method. J Ceram Int. 2015;41:7823-9.  7. Guo SQ. Densification of ZrB2‐based composites and their mechanical and physical properties: A review. J Eur Ceram Soc.  2009;29:995-1101.  8. Chamberlain AL, Fahrenholtz WG, Hilmas GE. Low‐temperature  densification of Zirconium Diboride Ceramics by reactive hot  pressing. J Am Ceram Soc. 2006;89:3638-45.  9. Chamberlain AL, Fahrenholtz WG, Hilmas GE. High‐Strength  Zirconium Diboride‐Based  Ceramics.  J Am  Ceram  Soc.  2004;87:1170-2.  F I G U R E 7   X‐ray diffraction data of specimen after oxidation  Figure 6A displays the SEM image of the cross‐section  perpendicular to the oxide layer. A continuous oxide layer  with a thickness of about 25\\xa0μm was observed on the surface  of the sample. This layer consisted of a mixture of SiO2,  Al2SiO5, and Al2O3 according to the EDS data of the surface and XRD data of specimen after oxidation (as shown in  Figures 6B and 7). When pure ZrB2 is oxidized at high temperature, the product B2O3 volatilizes leaving pores in ZrO2  on the surface of ZrB2. ZrB2 fur ther contacts oxygen and  gets oxidized. However, the oxidation resistance of ZrB2  could be improved by adding Al3BC3 and SiC. The oxide  products Al2O3 and SiO2 filled some pores generated by the  volatilization of B2O3. With the progress of the oxidation  reaction, a small amount of SiO2 and Al2O3 formed silicate  Al2SiO5. The mixture of Al2SiO5, SiO2, and Al2O3 formed a  continuous layer of oxides, which efficiently prevented oxygen from diffusing into the specimens dur ing oxidation, and  thus improved the oxidation resistance of ZrB2. Moreover,  this continuous layer of oxides repaired defects (such as processing damage and microcracks) on the surface of sample,  effectively reduced the size of defects, and reduce the loss  of mechanical proper ties to a cer tain extent.  4   |   C O N C L U S I O N S  ZrB2‐SiC‐Al3BC3 composite was successfully sintered by  HP at 1900°C. The microstructure of the sintered ZrB2‐SiC‐ Al3BC3 composite was dense and uniform. The densification  mechanisms were related to the liquid Al phase formed at low  temperature, reaction of carbon with the oxide impurities attached to ZrB2 particles, and the reaction between Al, B4C,  and C. The flexural strength and fracture toughness of the sintered ZrB2‐SiC‐Al3BC3 composite were 589\\xa0±\\xa0 147\\xa0MPa and   CHEN Et al.      \\x0c', '2196   |    10.   Jung JY, Jung SH, Oh HC, Lee SH, Choi SC. Spark plasma sintering of ZrB2 using Al3BC3 as an additive. J Ceram Process Res.  2012;13:641-5.  11. Fahrenholtz WG, Hilmas GE, Talmy IG, Zaykoski JA. Refractory  Diborides  of Zirconium  and Hafnium.  J Am Ceram Soc.  2007;90:1347-64.  12. Sha JJ, Li J, Wang SH, Zhang ZF, Zu YF, Flauder S, et al. Improved  microstructure and fracture properties of short carbon fiber‐toughened ZrB2‐based UHTC composites via colloidal process. Int J  Refract Met H. 2016;60:68-74.  13. Abdollahi A, Valefi Z, Ehsani N, Torabi S. High temperature anti‐ oxidation behavior of in situ and ex situ nanostructured C/SiC/ ZrB2‐SiC gradient coatings: thermodynamical evolution, microstructural characterization, and residual stress analysis. Int J Appl  Ceram Technol. 2018;15:1319-33.  14. Hwang SS, Vasiliev AL, Padture NP. Improved processing and  oxidation‐resistance of ZrB2 ultra‐high  temperature ceramics  containing SiC nanodispersoids. Mater Sci Eng A. 2007;464:  216-24.  15. Guo SQ. Textured microstructure and anisotropic strength of hot‐ forged ZrB2‐ZrC‐Zr cermets. J Ceram Int. 2016;42:16063-70.  16. Guo SQ. Densification, microstructure, elastic and mechanical  properties of reactive hot‐pressed ZrB2‐ZrC‐Zr cermets. J Eur  Ceram Soc. 2014;34:621-32.  17. Liu HT, Zou J, Ni DW, Wu WW, Kan YM, Zhang GJ. Textured  and platelet‐reinforced ZrB2‐based ultra‐high‐temperature ceramics. Scripta Mater. 2011;65:37-40.  18. Ni DW, Zhang GJ, Kan YM, Sakka Y. Highly textured ZrB2‐based  ultrahigh temperature ceramics via strong magnetic field alignment. Scripta Mater. 2009;60:615-8.  19. Akin  I, Hotta M, Sahin FC, Yucel O, Goller G, Goto T.  Microstructure and densification of ZrB2-SiC composites prepared  by spark plasma sintering. J Eur Ceram Soc. 2009;29:2379-85.  20. Wang HL, Feng L, Lee SH, Chen JB, Fan BB, Chen DL, et al.  ZrB2‐Al3BC3 composites prepared using Al‐B4C‐C additives and  spark plasma sintering. J Ceram Int. 2013;39:897-901.  21. Song JG, Li JG, Song JR, Zhang LM. Preparation of high‐density YAG/ZrB2 multi‐phase ceramics by spark plasma sintering. J  Ceram Process Res. 2007;8:356-8.   22. Xing XM, Chen JH, Bei GP, Li B, Chou KC, Hou XM. Synthesis  of Al4SiC4 powders via carbothermic reduction: reaction and grain  growth mechanisms. J Adv Ceram. 2017;6:351-9.  23 . Wang HL , Lee SH , Feng L . HfB2‐S iC compos i te prepared by reac t ive spark p lasma s in ter ing . J Ceram In t . 2014 ;40 :11009-13 .  24. Sciti D, Guicciardi S, Bellosi A. Effect of annealing treatments on  microstructure and mechanical properties of liquid‐phase‐sintered  silicon carbide. J Eur Ceram Soc. 2001;21:621-32.  25. Lee JK, Tanaka H, Kim H. Movement of liquid phase and the  formation of surface reaction  layer on  the sintering of β‐SiC  with an additive of yttrium aluminium garnet. J Mater Sci Lett.  1996;15:409-11.  26. Cao JJ, Moberlychan WJ, De Jonghe LC, Gilbert CJ, Ritchie RO.  In Situ Toughened Silicon Carbide with Al‐B‐C Additions. J Am  Ceram Soc. 1996;79:461-9.  27. Kim YW, Mitormo M, Emoto H, Lee JG. Effect of initial a‐phase  content on microstructure and mechanical properties of sintered  silicon carbide. J Am Ceram Soc. 1998;81:3136-40.  28. Yonenaga I. Thermo‐mechanical stability of wide‐bandgap semiconductors: high temperature hardness of SiC, AlN, GaN, ZnO and  ZnSe. Phys B: Condens Matter. 2011;308-310:1150-2.  29. Lee SH, Kim HD, Choi SC, Nishimura T, Lee JS, Tanaka H.  Chemical composition and microstructure of Al3BC3 prepared by  different densification methods. J Eur Ceram Soc. 2010;30:1015-20.  30. Kim JY, Kim YW, Mitomo M, Zhan GD, Lee JG. Microstructure  and mechanical properties of α‐Silicon Carbide Sintered with  Yttrium‐Aluminum Garnet  and  Silica.  J Am Ceram  Soc.  1999;82:441-4.  31. Guo WM, Zhang GJ. Oxidation resistance and strength retention of  ZrB2‐SiC ceramics. J Eur Ceram Soc. 2010;30:2387-95.  How to cite this article: Chen Z, Zhao X, Wang H,   et al. Preparation and properties of dense ZrB2  composite reinforced by elongated SiC and Al3BC3   grains. Int J Appl Ceram Technol. 2019;16:2190-  2196. https ://doi.org/10.1111/ijac.13265   CHEN Et al.    \\x0c']"
},{
  "_id": 211,
  "PDF": "Preparation, ablation behavior and mechanism of C-C-ZrC-SiC and C-C-SiC composites.pdf",
  "Text": "[\"Ceramics International 44 (2018) 7481-7490  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Preparation, ablation behavior and mechanism of C/C-ZrC-SiC and C/C-SiC composites  T  Zhigang Zhao, Kezhi Li⁎, Wei Li, Qing Liu, Gang Kou, Yulei Zhang⁎  State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi'an 710072, PR China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Carbon/carbon composites Precursor infiltration and pyrolysis Ablation Microstructure  ZrC precursor was synthesized by a solution approach using ZrOCl2·8H2O, acetylacetonate, glycerol and boronmodified phenolic resin. A ZrC yield of ~ 40.56 wt% was obtained at 1500 °C in the C/Zr molar ratio of 1:1. C/CZrC-SiC composites were fabricated by a combined processes of chemical vapor infiltration (CVI) and precursor infiltration and pyrolysis (PIP) using the synthesized ZrC precursor. For comparison, C/C-SiC composites were prepared by CVI. Thermogravimetric analysis showed that C/C-ZrC-SiC composites exhibited better oxidation resistance than C/C-SiC composites. After oxyacetylene torch ablation, the mass ablation rate of C/C-ZrC-SiC composites was 9.23% lower than that of C/C-SiC composites. The porous ZrO2 skeleton in the ablation center was prone to be peeled off by the flame flow, resulting in the higher linear ablation rate of C/C-ZrC-SiC composites. The oxide layers of ZrO2 and SiO2 were formed on the transition and brim region of C/C-ZrC-SiC composites and acted as effective heat and oxygen barriers. For C/C-SiC composites, the C-SiC matrix was severely depleted in the ablation center and the formed SiO2 layer in the brim region could protect the matrix against further ablation.  1.  Introduction  Carbon/carbon (C/C) composites as an excellent thermal structural materials have been widely used in aerospace components such as leading edges, nose tips and heat shields for reentry vehicles for advanced aircraft [1,2]. These components operate in severe conditions including higher temperatures, faster speeds, higher stresses, hostile environments and more [3]. Therefore, the improved ablation resistance of C/C composites are required to withstand the severe thermal conditions. Introducing ultra-refractory materials into C/C composites is an effective route to improve their ablation resistance. Ultra-high temperature ceramics (UHTCs), including the carbides, nitrides, and borides of some transition metals (TMs) such as Zr, Hf, Ti, Nb, and Ta, have received much attention in aerospace field [4]. These materials have the high melting points (> 3000 °C), good thermal-shock resistance, excellent ablation resistance and chemical erosion resistance [5]. Among the UHTCs, zirconium carbide (ZrC) exhibits unique properties such as high melting temperature [6], relatively low density [6] and excellent ablation resistance at high temperatures [7]. Moreover, the refractory ZrO2 with a melting point of 2770 °C which allows them to endure temperatures as high as 2500 °C [8]. Generally, ZrC-SiC ceramics have better high temperature performance due to the combination  of the passivating character of SiC and the high melting temperature of ZrC [9]. In addition, the molten SiO2-ZrO2 under ultrahigh temperature can seal the defects such as pores and cracks. Therefore, the anti-ablation performance of C/C composites can be improved by modifying the matrix with ZrC-SiC ceramics. Various methods, such as precursor infiltration and pyrolysis (PIP) [10-15], reactive melt infiltration (RMI) [16-20] or a combined processes [21,22] have been used to fabricate C/C-ZrC-SiC composites. PIP is a solution-based processing which can make ZrC and SiC homogeneous distribution into C/C composites. In this process, the low density C/C composites are selected as preform and pyrolytic carbon is commonly used as interphase layer to protect the carbon fiber from the chemical erosion by ZrC and SiC precursor during pyrolysis [14,15]. However, the significant difference of coefficient of thermal expansion (CTE) between carbon fibers and ZrC-SiC matrix will induce stress in the composites, which results in fiber-matrix interface cracking after heat treatment. SiC has a thermal expansion coefficient close to that of carbon matrix (SiC:3.8-5.12 × 10−6/K; PyC: 1-2 × 10−6/K) [11] and has the capacity to generate SiO2 layer under the oxidative conditions which acts as a diffusion barrier between oxygen and carbon surface [23,24]. Moreover, SiC layer can alleviate the CET difference existing between the fiber and matrix to a certain degree. So far, most literatures are focused on the ablation behavior of C/C composites only by  ⁎ Corresponding authors. E-mail addresses: likezhi@nwpu.edu.cn (K. Li), zhangyulei@nwpu.edu.cn (Y. Zhang).  https://doi.org/10.1016/j.ceramint.2018.01.125 Received 20 November 2017; Received in revised form 15 January 2018; Accepted 15 January 2018  0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c\", 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  modifying the matrix with ZrC and SiC. There are few reports on the ablation behavior and mechanism of C/C composites fabricated by modifying the matrix with ZrC and meanwhile protecting the carbon fiber using SiC transition interphase layer. In the present work, a low cost route was used to synthesize the liquid ZrC precursor for PIP. A combined processes of CVI and PIP were used to fabricate C/C-ZrC-SiC composites, in which the carbon fibers were protected by SiC coating and the matrix was modified using ZrC. C/C-SiC composites were fabricated by CVI for comparison. The phase and structure of ZrC precursor were characterized. Microstructure, oxidation resistance, ablation behavior and mechanism of the prepared composites were studied.  2. Experimental  2.1. Synthesis of ZrC precursor  Polyzirconoxane with a high stability and good reproducibility was chosen as the zirconium-containing precursor for the preparation of ZrC. Boron-modified phenolic resin was used as the major source of carbon due to its good solubility in polyzirconoxane and ethanol. The ZrC precursor was synthesized through a modified method similar to [25], and the schematic diagram is shown in Fig. 1. ZrOCl2·8H2O, acetylacetonate, glycerol and boron-modified phenolic resin were used as the starting materials in the synthesis of ZrC. The mass % yield of ZrOCl2·8H2O and boron-modified phenolic resin was determined using thermogravimetric analysis at the heating rate of 10 °C/min to 1500 °C under Ar atmosphere. The mass % yield for ZrO2 and C at 1500 °C was 41.71% and 8.86%, respectively. ZrOCl2·8H2O used as zirconium source was first dried at 60 °C for 3-5 h and then mixed with acetylacetonate (acacH) using molar ratio of 1 (acacH to Zr) and ethanol was used as a mutual diluent. Glycerol was added to the solution with the molar ratios in the range of 0.15-0.5 for glycerol/Zr under magnetic stirring. The resulting solution was refluxed at 140-180 °C for 2 h for the purpose of converting to polyzirconoxane-based solution. Appropriate amounts of the bulk boron-modified phenolic resin were added to the polyzirconoxane-based solution to give different C/Zr molar ratios (1:1, 1.2:1). The as-prepared ZrC precursors were dried at 60 °C for 10 h and cured at 180 °C for 3 h under air atmosphere for the purpose of evaporation of solvent and cross-linking of the precursors, respectively. The cured precursors were crushed to powder. Pyrolysis was performed in a graphite furnace for 2 h at 1500 °C at a heating rate of 2.5 °C/min under flowing argon. The ceramics yield was determined by the mass ratio of the pyrolyzed product at 1500 °C to the cross-linked precursor at 180 °C.  2.2. Preparation of materials  The preparation process of the C/C-ZrC-SiC and C/C-SiC composites is illustrated in Fig. 1. The density of the 2D needle-punched carbon felt was increased from 0.45 to 0.7 g/cm3 by depositing the pyrolytic carbon using isothermal chemical vapor infiltration (ICVI) at 900-1200 °C for 20 h, for the purpose of protecting the carbon fiber against the chemical erosion from the oxidative species produced by the ceramics precursor. The methane was used as carbon source. Then SiC was introduced into the preform of 0.7 g/cm3 by chemical vapor infiltration. The details of the process were described elsewhere [26]. The as-fabricated C/C-SiC composites were infiltrated with the liquid ZrC precursor by vacuum impregnation and pyrolysis for several times. Finally, the obtained C/C-ZrC-SiC composites were further densified with pyrolytic carbon using thermal gradient chemical vapor infiltration (TCVI) for 130 h. The same process was also employed to prepare the C/C-SiC composites. It should be noted that the final composites were not further processed at higher temperature. Mass changes of the composites before and after CVI (SiC) or PIP (ZrC) were recorded for evaluation of SiC or ZrC composition. The porosity of the composites was determined by Archimedes method. The bulk density of the composites was measured by dividing mass by volume. The phase compositions, density and open porosity of the asprepared composites are listed in Table 1.  2.3. Oxidation tests  2.3.1. Non-isothermal thermogravimetric analysis (TGA) The oxidation behavior of C/C-ZrC-SiC and C/C-SiC composites was studied by thermogravimetric analysis using a TA Q600 synchronous thermogravimetric analyzer. The bulk specimens with average weight of 32 mg were heated from room temperature up to 1400 °C at the heating rate of 20 °C/min in dry air with a flow rate of 30 mL/min.  2.3.2. Oxyacetylene torch testing The high temperature oxyacetylene torch tests were used to investigate the ablation resistance of the composites. The pressures and flow rates of C2H2 and O2 used to produce the torch with a heat flux of 4.18 MW/m2 were 0.095 MPa and 0.31 L/s, 0.4 MPa and 0.42 L/s, respectively. The inner diameter of the oxyacetylene gun tip was 2 mm. Plate-shaped specimen had a diameter of 30 mm and a thickness of 10 mm. The distance and angle between the specimen and the tip were 10 mm and 90°, respectively. The specimens were exposed to the torch for 60 s. Both the mass and linear ablation rates were calculated by the changes in mass and ablation center thickness as per unit time before and after test, respectively. The surface temperature of the specimen  Fig. 1. The schematic diagram for the preparation of  ZrC precursor, C/C-ZrC-SiC and C/C-SiC composites.  7482  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Table 1 Phase compositions, density and open porosity of the as-prepared composites.  Composites  Density (g/cm3)  Open porosity (%)  ZrC (wt%)  SiC (wt%)  C/C-ZrC-SiC C/C-SiC  1.93 1.99  6.32 6.85  14.46% -  28.68% 31.23%  during testing was measured using an infrared thermometer MR1SCSF).  (Raytek  2.4. Materials characterization  Fourier transform infrared spectra (FTIR, EQUINOX55) was used to investigate the structure of the ZrC precursors. Cross-sections of the fabricated composites were orderly ground flat using 800, 1000 and 1200 mesh SiC papers and then polished with diamond abrasives of 0.5 µm. X-ray diffraction (XRD, D/max-2500) with Cu Kα radiation was used to detect the phases for the pyrolyzed precursor powders and the as-prepared composites preand post-test. The microstructures of the samples were analyzed by a scanning electron microscope (SEM, VEGA3) combined with energy dispersive spectroscopy (EDS).  3. Results and discussion  3.1. FTIR of  the ZrC precursor  Fig. 2 shows FTIR spectra of the ZrC precursor without and with boron-modified phenolic resin additions. The spectrum for the pyrolyzed precursor powders without boron-modified phenolic resin additions is shown in Fig. 2a. The broad absorption peak at 3357 cm−1 is attributed to O-H stretching vibration. The strong bands at 1534 cm−1 (C˭C), 1350 cm−1 (CH3), 1287 cm−1 (C-H+C-CH3) and 1032 cm−1 (CH3) indicate the existence of acetylacetone ligand. The peak at 1454 cm−1 is attributed to C˭O stretching and C-H deformation [27]. The Zr-OC and ZrO-C stretching vibrations occur at 934 cm−1 and 1098 cm−1 [28]. The absorption peak at 652 cm−1 is due to Zr-O in the Zr-O-C chain, which indicates the Zr-O-C with the the formation of backbone of polymer in this new polymer [29]. The ZrC precursor with boron-modified phenolic resin additions (Fig. 2b) shows the formation of a broad absorption bands at 950 cm−1, which is due to B-O structure [30]. An absorption peak is observed at 751 cm−1 which is assigned to Zr-O-Zr chains. A new peak at 1103 cm−1, which attributes to the presence of B-H bonds [31]. The formation of bonding between ZrO2 and acetylacetone ligand and boron-modified phenolic resin groups may promote ZrC and/or ZrB2 formation during the pyrolysis process.  Fig. 3. XRD patterns of  the ZrC precursor powders  for different C/Zr molar  ratio,  (a)  without boron-modified phenolic  resin additions,  (b) 1:1 and (c) 1.2:1 pyrolyzed at  1500 °C for 2 h.  3.2. Phases of  the pyrolyzed precursor  XRD patterns for the ZrC precursors derived from different C/Zr molar ratios after pyrolysis at 1500 °C for 2 h are shown in Fig. 3. The pyrolyzed powders without the boron-modified phenolic resin additions are composed of m-ZrO2 and ZrC, shown in Fig. 3a. The formation of ZrC indicates that zirconium-containing groups are introduced into the organic molecular chain and the precursor achieves in situ carbonization during heating. The additional carbon source offered by boronmodified phenolic resin makes the C/Zr molar ratio to 1:1 (Fig. 3b) and the precursor has almost been transformed into ZrC by carbothermal reduction. In addition, a very small amount of ZrB2 and t-ZrO2 is formed. The ceramics yield is 40.56%. During heating, the chemical reactions in the precursor may occur to form ZrB2 according to reactions (1) and (2). When the carbon content increases and makes the C/ Zr molar ratio to 1.2:1, the pyrolyzed powders have the same phases as that of C/Zr molar ratio of 1:1 and the ceramics yield is 39.67% (Fig. 3c). High ceramics yield is a key factor to the efficiency of PIP process. Therefore, the ZrC precursor synthesized with C/Zr molar ratio of 1:1 is used for PIP.  ZrO2 + B2O3 + 5C = ZrB2 + 5CO (g)  ZrC + B2O3 + 2C = ZrB2 + 3CO (g)  (1)  (2)  3.3. Microstructures and phases of composites  the C/C-ZrC-SiC and C/C-SiC  The cross-sectional SEM images of C/C-ZrC-SiC and C/C-SiC composites are shown in Fig. 4. For the C/C-ZrC-SiC composites, the carbon fibers are coated by inner pyrolytic carbon layer and outer SiC layer with a thickness of 2.3 µm (Fig. 4a). The spaces between the individual carbon fibers are infiltrated with carbon matrix and white particles evidenced by EDS as ZrC-ZrO2 ceramics (Fig. 4c). XRD analysis shows that the C/C-ZrC-SiC composites are mainly composed of C, ZrC, SiC and a small amount of ZrO2 and ZrB2 (Fig. 5aI). A small amount of the formed ZrC is oxidized by oxidative species from the boron-modified phenolic resin during the repeated PIP process, resulting in the presence of ZrO2 in the matrix. The same results can also be found elsewhere [32]. As for the C/C-SiC composites, the microstructure of carbon fiber/ SiC coating is similar to that of C/C-ZrC-SiC composites (Fig. 4b) and the matrix is composed of carbon and SiC (Fig. 5aⅡ).  Fig. 2. FTIR spectra of the ZrC precursor powders cured at 180 °C for 3 h, (a) without the  boron-phenolic  resin additions and (b) with the boron-phenolic  resin additions  (C/Zr  molar ratio of 1:1).  Thermal analysis curves for the oxidation of C/C-ZrC-SiC and C/CSiC composites are shown in Fig. 6. The TGA curves show that no  3.4. Thermogravimetric analysis  7483  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Fig. 4. Microstructures of  cross  sections of  the as prepared composites  (a) C/C-ZrC-SiC,  (b) C/C-SiC;  (c) and (d) EDS analysis of  the marked area in (a).  evident weight change occurs up to 500 °C. C/C-ZrC-SiC and C/C-SiC composites experience a rapid weight loss in the temperature range of ~ 710 to 1195 °C and ~ 550 to 1204 °C, respectively. Weight loss is attributed to the oxidation of carbon. The DTG curves reveal that the maximum weight loss for C/C-ZrC-SiC and C/C-SiC composites occurs at about 1034 °C and 941 °C, respectively. TG-DTG shows that the oxidation of C/C-ZrC-SiC composites shifts to higher temperature compared with the C/C-SiC composites, indicating that C/C-ZrC-SiC composites exhibit better oxidation resistance relative to C/C-SiC composites. During oxidation, the diffusion barrier efficiency of the SiC coating on carbon fibers of the two composites is almost identical due to their similar microstructure of carbon fibers/SiC coating (Fig. 4a and b). Furthermore, SiC does not experience evident weight change in oxidative conditions from room temperature to 1400 °C [15]. The discrepancy in matrix composition of C/C-ZrC-SiC and C/C-SiC composites  causes their different oxidation properties. The carbon matrix of C/CSiC composites will be directly oxidized and oxygen diffusion through SiC layer is not the limiting process [23]. In this case, the oxidation rate is controlled by the reaction between carbon and oxygen. ZrC is more easily oxidized than C [15,33,34], which leads to a decrease in the consumption of carbon matrix. Thus, the oxidation resistance of C/CZrC-SiC composites is better than that of C/C-SiC composites.  3.5. Ablation behavior  3.5.1. Ablation properties of C/C-ZrC-SiC and C/C-SiC composites The ablation properties of the C/C-ZrC-SiC and C/C-SiC composites are listed in Table 2. Compared to C/ZrC-SiC, C/C-ZrC and C/C-SiC composites prepared using different methods elsewhere [35-37], both the C/C-ZrC-SiC and C/C-SiC composites in this study exhibit the better  Fig. 5. XRD patterns of C/C-ZrC-SiC (Ⅰ) and C/CSiC(Ⅱ) composites before (a) and after (b) exposure  to oxyacetylene torch for 60 s.  7484  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  SiO2 glass in gray is formed in the brim regions of both the C/C-ZrC-SiC and C/C-SiC composites samples. Different microstructures are formed in different regions of the samples after ablation. The ablated surface consists of SiC and C (Fig. 5b, Ⅱ), probably as a result of the visible pores on the surface exposing the underlying matrix. Microstructures of C/C-SiC composites in ablation center are shown in Fig. 8. It can be observed that SiC layer deposited on carbon fibers and C-SiC matrix are severely depleted by the flame. Needle-shape carbon fibers and shell matrix are formed (Fig. 8a and b) because of severe ablation of the carbon fibers parallel to the flame. EDS analysis reveals that the major elements of shell matrix are C, Si and O (Fig. 8e), implying that the shell matrix is composed of SiC and SiO2. Fig. 8(c and d) show the morphologies of the composites perpendicular to flame. The evident gaps between carbon fibers and matrix can be observed, indicating that the interface between SiC and carbon fibers is preferentially oxidized during ablation. Part of the shell matrix and fibers are even blown away by the flame and many honeycomb-like ablation etching pits generate in the remained matrix. Fig. 9 shows the microstructures of C/C-SiC composites from the edge region of the ablation center to the brim region. The matrix (mixture of SiC and SiO2) maintains well in the edge of the ablation center. Cracks can be found in the matrix because of the discrepancy in the CTE between SiC, SiO2 and C during cooling, as shown in Fig. 9a. The gap size between fiber and matrix is obviously decreased and tip dimension of the ablated carbon fibers is larger than that in ablation center. Therefore, the SiC layer offers an effective protection against ablation in this region where the temperature, flow velocity and pressure are lower. A mixture of SiC and SiO2 can be found on the carbon fiber surface according to the EDS analysis (Fig. 9d). SiO2 is formed after oxidation of SiC, which transforms into spherical particles driven by surface tension (Fig. 9b). Fig. 9c shows that a dense SiO2 layer is formed in the brim region of the sample. In this region, the temperature, pressure and flow rate of the flame flow are the lowest. The dense SiO2 layer can directly isolate the flow, restrain the oxygen diffusion, and hence prevent the composites against further ablation. The microstructures of C/C-ZrC-SiC composites after ablation are shown in Fig. 10. In the ablation center, a large amount of spherical particles can be observed on the top of carbon fibers parallel to the flame (Fig. 10a). Based on EDS analysis (Fig. 10f), these particles are mixture of oxides and carbides of Zr and Si [38]. SEM analysis reveals that there is no evidence of spreading of the spherical particles on the surface of carbon fiber (Fig. 10b). The particles adhere on the surface of carbon fiber and many pits can be found beneath these particles. Based on this case, it can be confirmed that the particles have poor compatibility with the carbon fiber. These particles are considered to exhibit two sides during ablation. On the one hand, they can prevent the flame from corroding the carbon fibers. On the other hand, they might react  Fig. 7. Digital  images of  samples after ablation (a)  C/C-SiC and (b) C/C-ZrC-SiC.  Fig. 6. TG-DTG curves for the oxidation of C/C-ZrC-SiC and C/C-SiC composites in air at  heating rate of 20 °C/min.  Table 2 Ablation properties of the carbon fiber composites fabricated by different process.  Composites  Process  Mass ablation rate (× 10−3 g/s)  Linear ablation rate (mm/s)  C/C-ZrC-SiC C/C-SiC C/ZrC-SiC C/C-ZrC C/SiC  PIP+CVI CVI PIP+RMI CVI+RMI PIP  2.95 3.25 13 -  27  0.015 0.009 0.022 0.028 0.064  Reference  This study  [35] [36] [37]  ablation resistance under high temperature conditions. The mass ablation rate of C/C-ZrC-SiC composites is 2.95 × 10−3 g/s, which is 9.23% lower than that of C/C-SiC composites. The linear ablation rate of C/C-ZrC-SiC composites is 66.7% higher than that of C/C-SiC composites. The poor mechanical denudation resistance of ablation products to the oxyacetylene torch is responsible for the high linear ablation rate of C/C-ZrC-SiC composites.  3.5.2. Morphology of the ablated surfaces Digital images of the sample surfaces after exposing to oxyacetylene torch for 60 s are shown in Fig. 7. As for the C/C-SiC sample, a large number of pores can be observed in the ablation center and the brim region of the sample remains structural integrity (Fig. 7a). For the C/CZrC-SiC sample, a visible spallation has occurred in the ablation center (Fig. 7b). The transition region of the sample is covered by a white oxide layer which is detected by XRD, shown in Fig. 5bⅠ, to be mostly ZrO2. The high speeds of the flow in the ablation center provide high shear stress which leads to the formation of the spallation. Whilst the  7485  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Fig. 8. Microstructures of the C/C-SiC composites in  the ablation center: (a, b) parallel to the flame, (c, d)  perpendicular to the flame and (e) EDS of the marked  A in (b).  with carbon fibers, resulting in many big oxidation etching pits left on the carbon fiber surface. However, the particles lie on the top of carbon fiber pits, indicating that their protection for carbon fiber is predominant. The SiC layer is depleted during ablation and carbon fibers perpendicular to the flame are severely ablated (Fig. 10c). A porous ZrO2 skeleton with big holes is formed in transition region (Fig. 10d). There are not enough molten ZrO2 to fill the holes and gaps between ZrO2 grains and CO might form at the ZrC-ZrO2 interfaces [39], which are the reasons for the formation of porous ZrO2 skeleton. In the brim region, the threadiness SiO2 with diameter of approximately 2 µm can be found because of oxidation of SiC (Fig. 10e). The growth mechanism of threadiness SiO2 is similar to that of silica nanowires without a metal catalyst via oxyacetylene torch ablation [40], which is ascribed to the growth catalyst provided by SiO2 particles and controlled by vapor-liquid-solid (VLS) mechanism.  3.5.3. Mechanism of ablation Surface temperature is one of quired for interpretation of test  the most important parameters reresults [41]. The time-temperature  curves of C/C-ZrC-SiC and C/C-SiC composites during ablation are presented in Fig. 11. It can be observed that the surface temperature rises up to 2150 °C at the rate of ~ 354 °C s−1 and to 1950 °C at the rate of ~ 332 °C s−1 during the initial ablation stage, and then increases slowly and maintains at 2340 °C and 2180 °C for C/C-ZrC-SiC and C/CSiC composites, respectively. It has been reported that formation of oxidation products leads to a change in the surface emissivity and thermal conductivity that affects the capacity to re-transmit heat [42,43]. Additionally, the matrix also has an important role in the heat conduction and thus affects the surface temperature. The thermal conductivity of SiC is higher than that of ZrC [6]. Increase in ZrC content should lead to decrease in thermal conductivity [44], indicating that C/C-SiC has a higher thermal conductivity in contrast to C/C-ZrCSiC composites. Compared with C/C-SiC composites, the higher surface temperature of C/C-ZrC-SiC composites is attributed to both the lower thermal conductivity of ZrO2 (~ 2.4 W (m K)−1) [45] and C-ZrC-SiC matrix. It should be noted that the surface temperature is dropped at 28 s, probably implying that ZrO2 in the ablation center is peeled off from the surface. As for the C/C-SiC composites, the evaporation of  7486  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Fig. 9. Microstructures of the C/C-SiC composites (a)  parallel and (b) perpendicular  to the flame in the  edge region of ablation center,  (c)  the brim region  and (d) EDS of the marked A in (c).  gaseous productions absorbs the energy from the flame and the C-SiC matrix is able to reduce the ablation heat of surface through conduction [46], resulting in a low surface temperature. Ablation is a complex physical and chemical reaction process which includes both high temperature oxidation and mechanical erosion from high enthalpy flows. During ablation, the C/C-ZrC-SiC and C/C-SiC composites are exposed to the oxygen rich environment. The major reactions in the composites are given below.  2C (s) + O2 (g) = 2CO (g)  C (s) + O2 (g) = CO2 (g)  C (s) + CO2 (g) = 2CO (g)  SiC (s) + 3/2O2 (g) = SiO2 (l) + CO (g)  SiC (s) + O2 (g) = SiO (g) + CO (g)  SiC (s) + 2SiO2 (l) = 3SiO (g) + CO (g)  SiC (s) = Si (g) + C (s)  SiO2 (l) = SiO2 (g)  SiO2 (l) + 3C (s) = SiC (s) + 2CO (g)  SiO2 (g) + 3C (s) = SiC (s) + 2CO (g)  ZrC (s) + 3/2O2 (g) = ZrO2 (s) + CO (g)  ZrB2 (s) + 5/2O2 (g) = ZrO2 (s) + B2O3 (l)  ZrO2 (s) + 3C (s) = ZrC (s) + 2CO (g)  ZrO2 (s) = ZrO2 (l)  B2O3 (l) = B2O3 (g)  (3)  (4)  (5)  (6)  (7)  (8)  (9)  (10)  (11)  (12)  (13)  (14)  (15)  (16)  (17)  In the case of C/C-SiC composite, the ablation rate is mainly controlled by the chemical reactions of carbon and SiC (reactions (3)-(10)). The interface between the matrix and carbon fibers is preferentially oxidized during ablation. The flame temperature in the ablation center is as high as 3000 °C. Part of SiC layer and SiC matrix are decomposed directly and undergo passive to active oxidation transition [6,47,48],  7487  resulting in the formation of SiO and CO instead of forming a protective SiO2 glasses layer. Thus, the matrix and fibers are severely oxidized and even broken away by the shearing force of oxyacetylene flame in the ablation center. Moreover, the reaction products, such as CO2, SiO, SiO2 (l) and SiO2 (g) produced by the oxidation of SiC and carbon will aggravate the erosion of the matrix and fibers (reactions (5), (8), (11) and (12)) (Fig. 9b, c and d). It can be found that the bare needle-shape fibers layer with many big gaps is formed on the surface owing to the severe depletion of matrix by the flame (Fig. 12a). In the brim region, the surface temperature of the specimen is in the temperature range of 1600-1700 °C [49] and the molten SiO2 layer is well retained (Fig. 12b). This SiO2 layer has ability to weaken the scouring of the oxygen acetylene gas flow and offer significant protection against ablation and isolate the penetration of oxygen into matrix. As for the C/C-ZrC-SiC composites, oxidation reactions between ZrC, SiC, C, ZrB2 and oxygen occur (reactions (3)-(17)). The major oxides are ZrO2, SiO2, SiO, CO, B2O3 and CO2. At the beginning of ablation, ZrC and ZrB2 are oxidized to ZrO2, leading to the high surface temperature. B2O3 evaporates rapidly from liquid due to its high vapor pressure at this ablation conditions. The severe depletion of C-SiC matrix and carbon fibers on the surface during ablation causes the formation of porous ZrO2 skeleton. The formed ZrO2 skeleton can withstand high temperature and play the thermal barrier effect to a certain degree. With continued ablation, the porous ZrO2 skeleton is prone to be peeled off by high-velocity gas flow due to poor configurational stability. The fresh matrix is exposed to the flame and ablated again. Oxidation of the C-SiC matrix leads to the formation of big gaps between the carbon fibers, which increases the thermal influence depth, as shown in Fig. 13a. In the transition region, oxygen passes through the porous ZrO2 layer and reacts with the underlying materials to form the inner layer of oxides and carbides of Zr and Si (Fig. 13b) based on EDS analysis in Fig. 13d. The temperature in the inner layer is believed to be lower than the surface temperature, and therefore the depletion of SiC is decreased. The formed inner layer can effectively prevent the further inward transport of oxygen. In the brim region, analysis by SEM-EDS of cross-section of the C/CZrC-SiC composites reveals that SiO2 as the outermost layer is formed on the surface (Fig. 13c and e). The similar structure can be found  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Fig. 10. Microstructures of  the C/C-ZrC-SiC compo sites after ablation:  (a)  the carbon fibers parallel  to  the flame in ablation center, (b) magnification of (a),  (c) perpendicular to the flame in ablation center, (d)  the transition region, (e) the brim region and (f) EDS  of  the marked A in (b).  (ZrO2+SiO2) is about 16.3% (corresponding to the marked A in Fig. 14) according to the ZrC and SiC content in the C/C-ZrC-SiC composites, assuming that ZrC and SiC are completely oxidized to ZrO2 and SiO2. It can be found that the minimum temperature of ZrSiO4 dissociation occurs at 1676 °C, above which t-ZrO2, cristobalite and liquid are formed. ZrSiO4 and SiO2 should be formed after cooling. However, the reverse reaction is suppressed during a high cooling rate, which results in the higher ZrO2 content and glass phase [52]. Especially, in the ablation conditions, the active oxidation of SiC and the decomposition of SiO2 form SiO which in turn reacts with oxygen to form SiO2 instead of ZrSO4 on the surface. Moreover, passive oxidation of part of SiC also contributes to the formation of SiO2 layer. Analysis by EDS of the matrix reveals that the presence of SiO2 layer is able to isolate oxygen diffusion into the matrix and resist mechanical denudation from the flame flow.  4. Conclusions  A economically efficient way was used to synthesize ZrC precursor  Fig. 11. The surface temperature curves of C/C-ZrC-SiC and C/C-SiC composites during  ablation.  elsewhere [50]. EDS analysis reveals that the matrix below the SiO2 layer mainly consists of ZrC, SiC and some oxides (Fig. 13f). From the eye of ZrO2-SiO2 phase diagram [51], in this study, the ratio of ZrO2/  7488  \\x0c', 'Z. Zhao et al.  Ceramics International 44 (2018) 7481-7490  Fig. 12. Cross-sections of C/C-SiC composites (a) the  ablation center and (b) the brim region.  which had a ceramics yield of ~ 40.56 wt% and can be transformed into ZrC at a relatively low temperature. C/C-ZrC-SiC composites were fabricated by a combined processes of CVI and PIP, the carbon fiber was first protected by SiC layer and the matrix was modified with ZrC. Thermogravimetric analysis showed that the C/C-ZrC-SiC composites exhibited a better oxidation resistance than C/C-SiC composites due to the preferential oxidation of ZrC which led to a decrease in the consumption of carbon. Ablation results showed that the mass ablation rate of C/C-ZrC-SiC composites was 9.23% lower than that of C/C-SiC composites. The porous ZrO2 skeleton was readily peeled off by the intense mechanical denudation of the flame flow in the ablation center, leading to the higher linear ablation rate of C/C-ZrC-SiC composites. The spherical particles (mixture of oxides and carbides of Zr and Si) formed on the carbon fibers were beneficial for slowing down or stopping the corrosion of carbon fibers by ablating flame. The ZrO2 and SiO2 layer were formed at the transition and brim region of C/C-ZrC-SiC and played the role of isolating the heat and restraining the oxygen diffusion to the underlying material. For C/C-SiC composites, the severe consumption of C-SiC matrix in the ablation center resulted in many big gaps formation and the SiO2 layer in the brim region could protect the matrix against further ablation.  Fig. 14. ZrO2-SiO2 phase diagram [51].  Fig. 13. Cross-sections and EDS analysis of C/C-ZrC-SiC composites at (a) the ablation center, (b) the transition region, (c) the brim region, (d), (e) and (f) EDS of the marked A, B and C,  respectively.  7489  \\x0c', 'Ceramics International 44 (2018) 7481-7490  [31]  [32]  [33]  [30]  [28]  [26]  [25]  [36]  [35]  [34]  [23] R. Gadiou, S. Serverin, P. 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},{
  "_id": 212,
  "PDF": "Preparation, microstructures, mechanical properties and oxidation resistance of SiBCN-ZrB 2 –ZrN ceramics by reactive hot pressing.pdf",
  "Text": "['Journal of the European Ceramic Society 35 (2015) 4399-4410  Contents lists available at www.sciencedirect.com  Journal  of  the  European  Ceramic  Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Preparation, microstructures, mechanical properties and oxidation resistance of SiBCN/ZrB2-ZrN ceramics by reactive hot pressing  Daxin Li, Zhihua Yang ∗ , Dechang Jia, Chengchuan Hu, Bin Liang, Yu Zhou  Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin 150001, China  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 5 June 2015 Received in revised form 5 August 2015 Accepted 10 August 2015 Available online 2 September 2015  Keywords:  SiBCN ceramics ZrB2 -ZrN ceramics Oxidation resistance Reactive hot pressing Mechanical properties  1.   Introduction  SiBCN/ZrB2 -ZrN ceramics were prepared by mechanical alloying followed by reactive hot pressing. The preparation, microstructures, mechanical properties and oxidation resistance were evaluated. Air exposure  immediate before cooling  leads to the formation of ZrB2 , ZrN, ZrO2 and SiC. Reactive hot pressing provides composite materials with similar structures consisting of ZrB2 , ZrN, SiC and a BNC phase. The mechanical properties of  the as-obtained ceramics depend on  the volume  fractions of  the  individual phases. High volume ratios of SiC/ZrB2 -ZrN in ceramics provide better oxidation resistance but reduce overall mechanical properties. The oxide surface of all investigated ceramics composes of SiO2 , ZrSiO4 and ZrO2 with a layered structure. The oxidation behavior of SiBCN/ZrB2 -ZrN ceramics in structure changes is similar, but the oxidation kinetics  is different. The oxide  layer can be divided  into three  layers: (1) outermost layer consists of SiO2 glass with uniform distribution of ZrSiO4, (2) porous ZrO2 -SiO2 layer, (3) unreacted matrix.  © 2015 Elsevier Ltd. All rights reserved.  The  transition-metal borides, carbides, and nitrides are classified as ultra-high-temperature ceramics  (UHTCs). UHTCs have recently attracted  significant  research  interest due  to potential applications  in thermal protection systems (TPSs)  for hypersonic aerospace  vehicles  and  reusable  atmosphere  re-entry  vehicles [1-3]. The UHTCs possess unique properties, including high melting points (>3000  C), excellent mechanical properties, strong oxidation resistance, and chemical  inertness  [4-6]. Because  it has  the lowest  theoretical density  (6.09 g/cm3 ) among UHTCs and good resistance to thermal shock result from its high thermal conductivity (65-135 W/m K), ZrB2 is an outstanding candidate for aerospace applications [3,7]. Numerous attempts have been made to enhance the oxidation resistance of ZrB2 -based materials using a wide variety of additives, especially SiC additives  [8-11], which  imparts oxidation resistance via  formation of dense SiO2 [12], as well as improved mechanical properties and sinterability [13]. The oxidation behavior and mechanism of SiC-ZrB2 have been well-defined [14-17]. When exposed  to air, ZrB2 oxidizes  forming crystalline ZrO2 and B2O3 . B2O3 is an effective oxygen diffusion barrier passivating the overall structure [18]. Unfortunately, B2O3 evaporates  ∗ Corresponding author at: Institute for Advanced Ceramics, School of Materials Science and Engineering, Harbin Institute of Technology, PO Box 3022, No. 2, YiKuang Street, Harbin 150080, PR China. Fax: +86 451 86414291. E-mail address: zhyang@hit.edu.cn (Z. Yang).  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.08.010 0955-2219/© 2015 Elsevier Ltd. All rights reserved.  dramatically above 1100  C reduces its effectiveness and coincidentally generating diffusion channels in the ZrO2 that remains thereby exposing the inner matrix to further oxidation [3]. At temperatures between 1100  C and 1400  C, para-linear oxidation kinetics have been observed. In this temperature regime, the rate of mass change is a combination of mass loss because of gas products and mass gain due to continuous oxidation [19]. Above 1400  C, it forms a layered structure consisting of: (1) a continuous silica-rich top layer, (2) a partially oxidized intermediate layer where SiC is depleted, (3) an unaffected bottom layer of SiC-ZrB2. Complimentary work on ZrN has  targeted  industrial  applications  including  hard  coatings,  diffusion  barriers,  optical applications  for  heat mirrors  and  decorative  coatings  [20]. ZrN has also been used as an  IR  reflective material  [21]. However, ZrN  is easily oxidized even at room  temperature  [22]. ZrN can be densified using pressureless sintering  (PLS), hot pressing (HP) or hot  isostatic pressing (HIP), generally requiring sintering temperature of up to 2000  C, long hold times and/or high applied loads due to  its high melting point, strong covalent bonding,  low self-diffusion coefficient and the presence of an oxide layer (ZrO2 ) on the ZrN powder surface [23]. Amorphous SiBCN materials are of great  interest as promising materials for aerospace and aviation applications due to their special  structures and high  temperature properties  [24]. These materials are normally amorphous at temperatures up to 1400  C with high oxidation and creep resistance as compared to monolithic SiC or Si3N4 [25-27]. Recently, mechanical alloying and hot pressing have been used  to prepare SiBCN amorphous powders                                              \\x0c', '4400   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  Table 1 Composition design and sintering parameters of SiBCN/ZrB2 -ZrN ceramics.  Samples   Powders   Mole ratio   Sintering parameters  SZBT1900  SZ1T1900  SZ2T1900  SZ3T1800  SZ3T1900  SZ3T2000   SZB  SZ1  SZ2  SZ3  SZ3  SZ3   Si   2  2  2  2  2  2   C   3  3  3  3  3  3   BN   1  1  1  1  1  1   Zr   0.5  0.5  1  2  2  2   B  1  1  2  4  4  4   1900   C/30 min/N2 1900   C/30 min/N2 1900   C/30 min/N2 1800   C/30 min/N2 1900   C/30 min/N2 2000   C/30 min/N2  Fig. 1. XRD patterns of MA Si-B-C-N-Zr amorphous powder:  (a) SZ3 powder prepared by one-step ball milling;  (b) SZ3 powder prepared by  two-step ball milling, inserted XRD pattern corresponding to Zr-B mixed powder); (c) SZ3 powder exposed to air  immediately after mechanical alloying; (d) SZ3 powder heated treatment at 1200   C for 60 min in N2 atmosphere.  and thereafter dense ceramics [28]. Mechanical alloying  is a  low cost, simple, synthesis method that generates sufficient mechanical energy,  fracture/shear,  localized high temperatures and  fusion to form amorphous powders [24]. Thus, it represents a useful method of preparing nanocrystalline materials at room temperature, especially recently metal carbides, borides and silicides [29,30]. Dense bulk SiBCN ceramics can be fabricated at 1900-2000  C under N2 (30 min) forming composites systems consisting of nanosized  ␤-SiC,  ␣-SiC and turbostratic BNC uniformly distributed  in an SiBCN matrix [31].  In the current effort we have explored the use of SiBCN as an additive in ZrB2 -ZrN to couple the properties of both systems as discussed above. We used the simplest synthesis method to produce the diborides, reaction of elemental powder. The ZrC-ZrN powder can be prepared by self-propagating hightemperature  synthesis  (SHS)  from mechanically activated Zr-C powder mixtures, which  is also known as combustion synthesis [23]. Self-propagating reactions generate temperatures sufficient to promote local Zr melting [32], which may decrease the sintering temperature. In current work, we adopt mechanical alloying  (MA) method to prepare Si-B-C-N-Zr amorphous powder and thereafter reactive hot pressing to densify SiBCN/ZrB2 -ZrN ceramics. To the best of our knowledge,  there are no  reports on processing ZrB2 -ZrN bulk  ceramics using  this approach. The present work employs SiBCN as additive  to prepare SiBCN/ZrB2 -ZrN ceramics  in order to enhance the oxidation resistance of ZrB2 -ZrN ceramics at temfrom 1100  C  to 1500  C. Below we begin by peratures  ranging  discussing, the preparation, microstructures and mechanical properties of SiBCN/ZrB2 -ZrN ceramics.  In particular, we focus on the oxidation kinetics and morphological changes  that occur during oxidation.  Fig. 2. Gibbs  free energy change versus temperature curves of the reactions predicted occurred during mechanical alloying and sintering process in N2 atmosphere.  2. Experimental procedures  2.1.   Processing  Commercial c-Si (99.2% pure, 45.0  \\u242em, Beijing Mountain Technical Development center, China), graphite  (99.5% pure, 8.7  \\u242em, Qingdao Huatai Lubricant Sealing S&T Co. Ltd., China), h-BN (98.0% \\u242em, Advanced Technology & Materials Co. Ltd., Beijing, pure, 0.6  \\u242em, Qingdao Huatai  China),  boron  (99.0%  pure,  2.3  Lubricant \\u242em, Sealing S&T Co. Ltd., China) and zirconium (99.5% pure, 72.8  Beijing Mountain Technical Development center, China) were used as raw materials. Mechanical alloying was performed  in a planetary miller (P4, Fritsch GmbH, Germany) with ball to powder mass ratio of 20:1. The composition design of Si:C:BN:B:Zr  in mole ratio and sintering parameters are presented in Table 1 according to our previous research  [24]. The  rotation  speed of  the main disk was  set at 350 rpm/min and the vials were 600 rpm/min in reverse. The powder mixture was  loaded  into  the silicon nitride vials along with silicon nitride balls under Ar. The machine stopped  for 20 min  in every 40 min, and the total effective ball milling time was in a range of 1-40 h. The amorphous Si-B-C-N-Zr powders at target compositions were fabricated by oneand two-step mechanical alloying methods. For one-step milling, c-Si, h-BN, graphite, boron and zirconium were simultaneously milled for 20 h, denoted as SZ1, SZ2 and SZ3 powders according to Zr/B composition. As  for two-step milling, c-Si, graphite and h-BN powders in mole ratio of 2:3:1 were milled  for 20 h, and then the as-milled Zr-B powders  (5 h) were added and milling continued for another 5 h giving SZB powder. As-milled mixed powder was  load  into a graphite die (diameter = 50 mm) lined with graphite foil and coated with BN and then  \\x0c', 'D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410   4401  Fig. 3. Surface morphologies of  the as-milled Si-B-C-N-Zr amorphous powder before and after exposing to air: (a) SZ1 powder prepared by one-step ball milling; (b) SZ1 powder exposed to air instantaneously before cooling; (c) SZB powder prepared by two-step ball milling; (d) SZB powder exposed to air instantaneously before cooling. Inserted are the corresponded particle size distributions.  ×  applied  to reactive hot pressing  (Highmulti-10000). The  furnace was heated to 1200  C at an average rate of 25  C/min, then further heated to target temperatures (1800, 1900 or 2000  C) at 20  C/min. The  furnace was held at the target temperature  for 30 min  in N2 (1 bar) and then cooling to room temperature in an irregular way. load was removed at 1800  C after sintering. The The uniaxial  as-obtained monoliths were polished and  then cut  into bars of 36   4 mm  (30 mm outer space)  to measure flexural strength and Young’s modulus at room  temperature by  three point bend testing with a crosshead speed of 0.5 mm/min. The fracture toughness was measureed using  single  edge notched beam  (SENB) method of 20   4 mm  (16 mm outer  space) using a universal testing machine (Instron-5569, USA). Vickers hardness (HVS-5, Laizhou, Huayin, Testing Instrument Corp., USA) was measured on polished surfaces with a load of 10 kg and a hold time of 10 s. Five samples were used for each mechanical properties measurement and  the results are reported as an average. Oxidation  tests used 10   4 mm bar diced from the hot pressed billets.  ×  ×  ×   3   ×  ×   3    2   2.2. Oxidation  Oxidation  tests were  run  using  a  horizontal  tube  furnace (RHTH120/600/18, Nabertherm, Germany) equipped with a MoSi2 heating element. Before each test, the ceramics were polished with a diamond abrasive, down to a 0.5  \\u242em finish. To avoid the influence of mass change brought by surface moisture, ceramics were placed into an oven preset at 120  C for 10 h. Oxidation experiments were conducted in flowing air (N2 79%, O2 21%, steam content <50 ppm) at 1100, 1200, 1300 and 1500  C, respectively. The cleaned ceramics were placed in an alumina boat using alumina bars as support points, and then  inserted  into center of the  furnace and  leveled. Ceramics were heated to target temperatures at 5  C/ min and held for times ranging from 1-5 h.  2.3.   Characterization  To analyze  the microstructures of as-prepared powders and ceramics, a scanning electron microscope (SEM, 30 kV, Quanta 200 FEG, FEI Co., USA) and a transmission electron microscope (TEM,  Fig. 4. TEM and HRTEM images of SZ1 powder before and after exposing to air: (a), (b) SZ1 powder prepared by one-step ball milling after cooling; (c), (d) SZ1 powder exposed to air instantaneously before cooling. Inserted are the corresponded SEAD patterns and EDS result.  Table 2 Mechanical properties and densities of SiBCN/ZrB2 -ZrN ceramics.  Samples   Density (g/cm3 )  Flexural strength (MPa)  SZBT1900  SZ1T1900  SZ2T1900  SZ3T1800  SZ3T1900  SZ3T2000   3.22  3.24  3.65  3.96  3.97  4.08   213.0  202.0  226.0  222.1  302.9  400.0   ± ± ± ± ± ±   7.8   9.1   4.2   9.1   1.7   15.4   Fracture toughness (MPa m1/2 )  2.57  2.34  2.82  2.83  2.91  3.16   ± ± ± ± ± ±   0.41   0.24   0.31   0.34   0.24   0.19   Elastic modulus (GPa)  252.4  242.1  214.5  142.1  197.3  251.6   ± ± ± ± ± ±   5.8   4.6   10.6   4.6   6.9   5.8   Vicker’s hardness (GPa)  3.29  3.56  5.24  6.10  6.01  9.57   ± ± ± ± ± ±   0.57  0.88  1.24  0.53  1.86  1.24  Tecnai.F30, 300 kV, FEI Company., USA) were used. Energy dispersive spectroscopy (EDS, Oxford instruments INCAx-act, Oxforshire, U.K.) was employed  for chemical analysis. The structural characterization was analyzed by X-ray diffraction (XRD, 40 kV/100 mA, D/max-␥B Cuk␣, Rigaku Corp.,  Japan) method  to obtain X-ray diffraction patterns with a scanning speed of 4  /min.  3. Results and discussion  3.1.  Preparation, microstructures and mechanical properties of as-prepared powders and ceramics  Fig. 1(a) illustrates the XRD results of SZ3 powder prepared by one-step ball milling process. It is evident that there is no reaction between Zr, B, BN and C particles after ball milling for 1 h since Zr and Si phases are still observed  in the powder blends, and only a small amount of SiC  is detected due to reaction of Si and C. The absence of peaks of other components such as boron, graphite and h-BN are due to their amorphous states. However, it is found that the peaks of raw materials decrease in intensity dramatically and broaden with  increasing milling  time duration. When  the starting powders are ball milled for 3 h, ZrB2 is obtained and no traces of other compounds are observed. These results present the size reduction and the increase in the lattice strain of Zr and Si metal particles, and the disappearance of the stacking order of the graphite and the formation of disordered carbon with fine grains.  \\x0c', '4402   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  Fig. 6. Surface and  fracture morphologies of  (a),  (d) SZ3T1800 ceramics;  (b),  (e) SZ3T1900 ceramics; (c), (f) SZ3T2000 ceramics. (The white and gray areas observed on polished surface  indicate  the existence of ZrB2 -ZrN and SiBCN agglomerates, respectively).  SiBCN powder pre-milled  for 20 h. As  seen  in Fig. 1(b), all  the as-prepared powders are  in well amorphous states after milling for 20 h and more.  It  is known that ball milling parameters, such as rotation speed of vials, ball  to powder mass rate and milling time, may  influence the powder characteristics greatly, which  in turn may have a strong impact on the compressibility, sinterability and microstructures of  the powder compact  [24]. According  to literature [32], dense ZrB2 -ZrN ceramics can be synthesized using Zr and h-BN powders as  raw materials  combining mechanical alloying and high-pressure high-temperature techniques. Results showed that ZrN was observed after ball milling raw materials for 10 h  in a mole rate of 3:2. However,  in current work, ZrN  is not found but SiC, ZrB2 are obtained during ball milling. The reasons for this will be discussed in next section. Fig. 1(c)  indicates the XRD pattern of SZ3 powder exposed to air  immediately before cooling down. SZ3 powder exposed to air instantaneously enables complete solid-state chemical reactions of raw materials with oxygen, and the formation of ZrN, ZrB2 , SiC and ZrO2 are confirmed. Self-propagating high-temperature synthesis (SHS) which is also called combustion synthesis could lead itself to preparation of diboride-containing composites or diborides doped with additives such as sintering additives [4]. Tsuchida [33] prepared ZrC and ZrB2 in air  from Zr/B/C powder mixture  in a mole rate of 1:1:1 by mechanical activation assisted self-propagating high-temperature synthesis (MA-SHS) method. They found that the self-propagating high-temperature synthesis was because of  the reactions of disordered carbon formed by grinding with oxygen in air, followed by the reaction between Zr and disordered carbon as MA-SHS process. However, no trace of ZrN was detected  in their  Fig. 5. XRD patterns of SiBCN/ZrB2 -ZrN ceramics prepared by MA plus RHP methods.  Continuous deformation of powder particles through prolonged milling time could result  in crystallite refinement and  increase  in lattice strain, as indicated in Fig. 1(a). After ball milling for 20 h, the SZ3 powder can be regarded as completely amorphous structure. Since prolonged milling  time  cannot  change  the powder  input into powder,  the microstructures and  the crystallite content  in blends may have little variation despite ball milling time reaching to 40 h. For  this reason, 20 h ball milling  time  is adopted  in  the following process. Fig. 1(b) shows the XRD patterns of SZ3 powder prepared by  two-step ball milling method  at different milling time, and  inserted are  the XRD  spectra of Zr-B mixed powder pre-milled. With milling  time  reaching  to 1 h, ZrB2 forms and a small quantity of Zr  is  remained  in mixed powder. However, amorphous powder  is obtained after ball milling  for 5 h. Subsequently,  the Zr-B amorphous powder  is added  to  jar containing  \\x0c', 'D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410   4403  Fig. 7. Surface and fracture morphologies of the as-prepared ceramics: (a), (b) SZBT1900 ceramics; (c), (d) SZBT 2000 ceramics.  study.  In order to  identify whether or not there exists remained Zr particles  in SZ3 powder after MA, hence SZ3 powder  is heated treatment at 1200  C for 60 min in N2 , as presented in Fig. 1(d). The result suggests that ZrN, ZrB2 , SiC and a small amount of ZrO2 are also  formed  in N2 atmosphere, which  is consistent with Fig. 1(c) result. It seems that the formation of ZrN results from the reaction between residual Zr and N2 , besides that, the disordered carbon reacted with oxygen ignites the SHS process. The formation of ZrO2 in N2 atmosphere may come from the impurity of N2 used, which contains little amount of oxygen. In order to clarify the formation of ZrB2 , ZrN, ZrO2 and SiC, the Factsage software  is used to calculate the reaction energy. Gibbs free energy change versus  temperature curves of  the predicted occurred reactions of ZrB2 , ZrN, ZrO2, ZrC, B4C, ZrSi2 and SiC are displayed in Fig. 2 [30]. As we can see, the Gibbs free energy changes of all the predicted reactions are negative which mean that they may take place and therefore they are thermodynamically feasible at temperatures between 300 and 2000 K. Moreover, all the reactions shown above are exothermic which  imply  that an amount of heat release during mechanical alloying or sintering procedure. The formation of ZrB2 and ZrN rather than ZrC and B4C are duo to their higher Gibbs free energy change than others,  indicating the formation of ZrB2 and ZrN are more preferential than ZrC and B4C. According to Tsuchida’s work [23], the self-ignition reactions are due to the oxidation of disordered carbon formed by MA with oxygen in air. This is because the reactivity of the nanosized particles, such as disordered carbon, is extremely high. When the amorphous powders with high surface free energy contact with oxygen before cooling, the disordered carbon reacts with oxygen to release heat first. Despite the oxidation of Zr particles in air are expected to occur preferentially based on the thermodynamic date, but Zr particles are enwrapped with finely divided particle of disordered carbon and amorphous boron after MA process, therefore its oxidation process must be retarded. For homogeneously amorphous powder  is obtained by MA, the heat released by oxidation of disordered car bon is transferred to the nearby powder effectively. Hence SHS process in Si-B-C-N-Zr powder is induced. For both formation reactions shown in Fig. 2 are exothermic, so that the reactions become self-sustaining and propagate  through overall  reactant mixture. However, the formation of SiC in SZ3 powder after exposing to air may result from the formation of SiC nanograins during MA process and the absence of ZrC  is due to exhaustion of disordered carbon and Zr which remains further study. In summary, mechanical alloying can be employed to prepare amorphous Si-B-C-N-Zr powder incorporating a little amount of ZrO2 with fine microstructure. The exposure of mechanical activation disordered carbon to air results in an ignition reaction between carbon and oxygen, therefore MASHS process can be  induced and sustains. The  formation of ZrN may due to the reaction of residual zirconium and N2 . In general, the mechanical activation enhances the reactivity of solids as well as  their mixing uniformity, and SHS process can release a great amount of heat during exothermic reactions taken place. The coeffect shown above would  lower  the sintering  temperature and enhance the sinterability remarkably during reactive hot pressing. Fig. 3 shows the morphologies of as-milled SZ1 and SZB powders before and after exposing to air. After ball milling for 20 h, SZ3 and SZB powders mainly consist of near-spherical agglomerates, whose diameters are below 500 nm. The sizes are in well agreement with the particle size distribution collected by scattering particle-size analyzer, as indicated in Fig. 3(a) and (c). After exposure of SZ1 and SZB powders to air instantaneously before cooling, some particles seem to react with each other and form larger agglomerates. Trace of oxide is observed on powder surface, indicating ZrO2 formed in blends, which is matched with the XRD results. Fig. 4 displays the TEM and HRTEM  images of the SZ1 powder before and after exposing to air. As  indicated  in Fig. 4(a) and (b), the as-milled powder has amorphous structure which can be confirmed by two pieces of evidence: the completely disordered atomic arrangement in HRTEM image and corresponded SEAD pattern with large diffraction spot. However, the exposure  leads to the forma \\x0c', '4404   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  ±  ±  ±  ±  ±  ±  The mechanical properties and densities of SiBCN/ZrB2 -ZrN ceramics are  shown  in Fig. 8, and  the  related values are  summarized  in Table 2. For ZrN phase  forms  from  residual Zr and N2, we  could not use  the volumetric  rule of mixtures  to  calculate  the  theoretical density  for SiBCN/ZrB2 -ZrN  ceramics. As indicated,  the  fracture  toughness, flexural strength and Vicker’s hardness maintain  a  rising  trend, while  the Young’s modulus declines first but  increases rapidly at  last.  It can be seen that the sintering  temperature has a great  impact on mechanical properties due  to higher density obtained by higher  temperature. With the increasing content of Zr-B mixture, both fracture toughness and Vicker’s hardness  improve  steadily while  the flexural strength  remains unchanged. Meantime,  the Young’s modulus drops obviously from 242.1   4.6 GPa to 142.1   4.6 GPa. However, it  seems  that  the mechanical properties of SZBT1900  ceramics prepared by two-step ball milling plus RHP methods perform better than SZ1T1900 ceramics although they have almost the same density. The SZ3T2000 ceramics performs superior on mechanical properties and density compared with others, the flexural strength, fracture toughness, Young’s modulus and Vicker’s hardness reach 0.19 MPa m1/2 ,  ing  to  400.0   15.4 MPa,  3.16  251.6   5.8 GPa and 9.57   1.24 GPa, respectively. According to  literature [34-36], the Young’s modulus of ZrB2 -based ceramics with and without 350 GPa  sintering  additives  ranges  from  to  530 GP, depending on porosity and  sintering additives. The  fracture  toughness and flexural  strength are  in  the  range of 2.4-4.8 MPa m1/2 and 350-580 MPa,  respectively, depending  on  grain  size,  additives and  relative density  [37-41]. To obtain a  full densification and advantageous performance of UHTCs, fine grain size of starting materials and homogeneous microstructure are required. Reactive hot pressing  is a  set of  techniques  that utilizes high  temperature  reactions between  the  starting powders  to  simultaneously enhance densification, reduce grain growth, and potentially incorporate additional phases  [42]. Compared with conventional hot pressing, reaction sintering could reduce  the sintering  temperature as well as soak time required to prepared full dense ceramics. However,  the  research of  reaction  sintering  to prepare dense ZrB2 -based ceramics are limited and still remain room for further investigation compared with other methods, such as hot pressing, pressureless sintering and spark plasma sintering. Rosenberger [42] prepared Zr1−x TixB2 -ZrC composites with desirable microstructure and mechanical properties from ZrB2 -TiC powders by reactive hot pressing. Results showed that the flexural strengths and hardnesses of the ceramic composites sintered with TiC were greater than the conventionally processed ZrB2 -ZrC materials, rising from 440 MPa and 17.4 GPa to a maximum of 670 MPa and 24.2 GPa at 10 vol% TiC. In a similar work by Wang [43], they prepared ZrC-SiC composites 1600  C with relative densities  at  low temperature  in excess of 98% using ZrC and Si powders by reactive hot pressing. They also reported an increase of fracture toughness and hardness compared with pure ZrC ceramics. Besides, Guo [44] prepared ZrB2 -ZrCx -Zr cermets with 20 vol% SiC additives via  reactive hot pressing of Zr + B4C + SiC powder mixtures and obtained the flexural strength of cements in the range of 384.2-409.2 MPa. However, the mechanical properties presented in current work are lower than the literature mentioned above, this should attribute to the volume fractions of the  individual phases.  In our previous work [45], nanocrystalline SiBCN  ceramics were prepared  from  the mechanically-alloyed amorphous SiBCN powder, hot pressed at 1900  C under 80 MPa in the nitrogen atmosphere for 30 min. The as-obtained ceramics have room-temperature density, flexural strength, Young’s modulus and fracture toughness of 2.6 g/cm3 , 331.1 MPa, 139.4 GPa and 2.8 MPa   m1/2 , respectively.  Incorporation of SiBCN component  in ZrB2 -ZrN matrix may result in decreasing the mechanical properties of ceramic composites compared with pure ZrB2 -ZrN ceramics.  Fig. 8. Mechanical properties of SiBCN/ZrB2 -ZrN ceramics: (a) flexural strength; (b) fracture toughness; (c) elastic modulus; (d) Vicker’s hardness.  tion of nanosized grains at  local region, seen  in Fig. 4(d). The EDS result shows the nanocrystalline region consisting of Zr, Si, B, C, N and O elements, which is consistent with the XRD results. Fig. 5 presents the XRD patterns of SiBCN/ZrB2 -ZrN ceramics prepared by reactive hot pressing. After sintering, obviously crystallized peaks are observed and the main phases are those of ZrB2 , ZrN, SiC, a small amount of BNC, ZrO2 and c,  t-ZrO2 (marked as ZrOx ). The diffraction peaks of 33.97  , 39.42  , 56.75  and 71.42  correspond to (1 1 1), (2 0 0), (2 2 0) and (3 1 1)  lattice planes of cZrN, while peaks of 25.8  , 32.96  and 42.10  correspond to (0 0 1), (1 0 0) and  (1 0 1)  lattice planes of h-ZrB2 [32]. Besides, peaks of 28.31  and 31.49  correspond to ( ¯1 1 1) and (1 1 1) lattice planes of m-ZrO2, while peaks of 31.13  is matched with c, t-ZrO2 . However, it is interesting that the formation of ZrN and ZrB2 occurs predominantly and a trace of zirconium oxide is also detected. These results indicate that the residual Zr in amorphous powders prefers to react with N2 duo to its higher Gibbs free energy change than others during sintering. Furthermore, the volume ratios of SiC/ZrB2 -ZrN and ball milling methods (one-step or two-step milling) have an impact on the diffraction  intensities. With a higher content of Zr-B mixture, the diffraction peaks of ZrB2 and ZrN increase along with the decrease intensity of SiC, seen in Fig. 5(b). As presented in Fig. 5(c), the main diffraction peaks of SZBT1900 ceramics are higher than that of SZ1T1900  ceramics,  indicating a good  crystallization of SZBT1900 ceramics. The microstructure  representatives  of  the  SiBCN/ZrB2 -ZrN ceramics are shown  in Fig. 6. SEM  images present the  indication of macro-porosity and cracks distributed around at relatively low sintering temperature, however, a dense ceramics can be obtained after rising the sintering temperature which is consistent with the density values measured by the Archimedes method, as displayed in Table 2. Base on SEM observations and density measurements, porosity is not expected to impact the oxidation behavior. On polished surfaces, the gray areas are SiBCN gathering agglomeration, while the white ones indicate the ZrB2 -ZrN component. From the fracture morphologies of as-prepared ceramics, almost no pores can be  found and ZrB2 -ZrN  shows  isometric grain morphology while SiBCN exhibits laminated structure. Fig. 7 shows the surface and  fracture morphologies of SZBT1900 and SZBT2000 ceramics. Again, no obvious pores and cracks are found on surface and fracture morphologies. It seems that the impact of ball milling method on microstructure of the as-prepared ceramics remains small and shows a  similar  structure, besides,  the grains of ZrB2 -ZrN are homogenously distributed with average size  less  than 1  for both ceramics.  \\u242em   ·  \\x0c', 'D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410   4405  in matrix, the more ZrSiO4 exists in amorphous SiO2 glass. Beyond that, no distinctive feather is discovered. Fig. 11 shows the surface and cross-sectional morphologies of SBZT2000 ceramics oxidized at 1500  C for 5 h. Fig. 11(a) exhibits a distinct dividing line—a boundary between the porous region and glassy zone. The enlarged image of dense area shown  in Fig. 11(c) reveals that the uniform distribution of ZrSiO4 among  the  fused SiO2 layer have a size below 20  \\u242em. The corresponded EDS  results confirm  the above analyses, which are matched with the XRD result. Fig. 11(b) displays the cross-section of SZBT2000 ceramics after oxidation. A full and dense oxide  layer with relatively uniform thickness adhered to ceramic matrix  is obtained. Furthermore, the oxide  layer  is about 220  \\u242em in thickness without any penetrating cracks or holes between the oxide  layer and the matrix. Similar to ZrB2 -ZrC ceramics [12,14], a  layered structure of oxide scale  is  formed during oxidation of SiBCN/ZrB2 -ZrN ceramics at 1100-1500  C. In general, two methods are measured to assess the oxidation kinetics of ceramic composites: weight change and thickness measurement. Structure ceramics, such as ZrB2 -SiC and ZrB2 -MoSi2 , are used weight variation evaluation to describe oxidation kinetics [33,46]. Gibbs free energy and reaction enthalpy changes of the following reactions predicted occurred on oxide  layer are shown in Fig. 12. Besides  that,  the phase diagrams of B2O3 -SiO2 and ZrO2 -SiO2 would be helpful  to  illustrate  the oxidation structure as well as kinetics.  →  →  →  +  +  +  +  +  +  +  +  4BN(s)   2ZrN(s)    3O2 (g)    2B2O3 (l)    2N2 (g)    2O2 (g)    2ZrO2 (l)    N2 (g)   2ZrB2 (s)    5O2 (g)    2ZrO2 (l)    2B2O3 (l)   SiC(s)    O2 (g)    SiO(l)    CO(g)   →  +  2SiC(s)    3O2 (g)   →   2SiO2 (l)   +   2CO(g)   (1)  (2)  (3)  (4)  (5)  Fig. 9. XRD patterns of SiBCN/ZrB2 -ZrN ceramics prepared by reactive hot pressing after oxidation test: (a) SZ3T2000 ceramics oxidized at different temperatures for 5 h; (b) ceramics oxidized at 1500   C for 5 h.  3.2. Oxidation resistance of SiBCN/ZrB2 -ZrN ceramics  The XRD spectra of resulting ceramics after oxidation test are shown  in Fig. 9. ZrSiO4 , SiO2 and ZrO2 phases are observed, however, no trace of boron oxide is detected. The intensities of ZrSiO4 in all investigated ceramics are much higher than others, showing a main oxidation product of ZrSiO4 on oxide surface. The diffraction  intensities of SiO2 and ZrO2 phases are not evident at all oxide surfaces. The diffraction intensities corresponding to ZrSiO4 increase rapidly at elevated oxidation temperature, suggesting an enhancement of thickness of oxide  layer as displayed  in Fig. 9(a). Compared with SZBT1900 ceramics, the diffraction  intensities of ZrSiO4 on SZ1T1900 surface are stronger than the  former, which indicates that the mechanical alloying process (one-step or twostep) can influence the oxidation behavior. The lower volume ratios of SiC/ZrB2 -ZrN, the higher amount of ZrSiO4 can be obtained, as exhibited in Fig. 9(b). SEM  investigations of oxide surface of  the resulting ceramics oxidized at 1100  C  for 5 h are shown  in Fig. 10. No visible pores and cracks are found on oxide surface except Fig. 10(a). As indicated, the oxide surface of SZBT1900, SZ1T1900, SZ2T1900 and SZ3T1800 ceramics are covered with a SiO2 glass  layer which appears  to have been liquid and amorphous at the oxidation temperature, and the EDS results  inserted also confirm the existence of amorphous SiO2 covered on surface. For SZBT1900 ceramics, the pores remain on surface resulting from the evaporation of gas products. For all the ceramics prepared by one-step ball milling powder,  the  formation of ZrSiO4 exists  in amorphous SiO2 which has an  island structure. The higher content of ZrB2 -ZrN component distributes  The above reactions are both extremely exothermic and favorable at all target temperatures from 300-2000 K, as shown in Fig. 12(b). It is well known that below 1100  C, the oxidation rate of SiC is much slower than ZrN and ZrB2 , and that above 1100  C, SiC is oxidized to form SiO2 rapidly [14]. The oxidation driving force of ZrB2 is the highest than others, followed by SiC, BN and ZrN  in turn. According to Lu’s work [47], the formation temperature of B2O3 is below 1100  C during oxidation test despite its Gibbs free energy is lower than others. The formed B2O3 spread on oxide surface and evaporate dramatically at elevated temperature, meantime SiO2 cover and spread on surface and  form oxide  layer gradually. However, SiO2 and B2O3 could  react with each other  to  form borosilicate glass in terms of SiO2 -B2O3 phase diagram. Furthermore, the oxygen  threads  the existence of borosilicate  layer easily. For  those reasons,  the  formation of borosilicate  layer would be destroyed completely due to the separation of SiO2 and volatilization of B2O3 . As a result, the borosilicate  layer disappears at this temperature regime, which  is consistent with XRD and EDS results. The catastrophic destruction of borosilicate layer reduces the effectiveness of the diffusion barrier, therefore could not protect the inner matrix from being oxidized. Meanwhile ZrB2 , ZrN and SiC are oxidized to form SiO2 and ZrO2 accompanied with gases evaporation, thus pores and cracks are left behind on oxide layer. For the porous structure of ZrO2 left behind does not retard the oxygen into underlying matrix effectively, a continuous and relatively high oxidation rate would emerge. Nevertheless, the formation of substantial SiO2 covered on oxide surface heals pore and cracks sufficiently despite its high viscosity at  low oxidation temperature, as shown  in Fig. 10. Tang [48] employed polyborosilazane-ethanol solution as precursor and a  few ZrB2 particles  to prepare a dense and  seamless ZrSiO4 /SiBCN(O) amorphous coating at 1500  C. They found that the borosilicate glass still exist on oxide layer in medium temperature of 1273-1473 K, which had a low viscosity and high wettability. At  \\x0c', '4406   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  Fig. 10. Surface mophologies of the resulting ceramics oxidized at 1100   C for 5 h: (a) SZBT1900 ceramics; (b) SZ1T1900 ceramics; (c) SZ2T1900 ceramics; (d) SZ3T1800 ceramics; (e) SZ3T1900 ceramics; (f) SZ3T2000 ceramics.  high temperature of 147--1773 K, the high-thermally stable ZrSiO4 phase formed by the reaction between ZrO2 and SiO2 , which could improve the stability of silica glass and lower the diffusion rate of oxygen [49]. However, in our current research, ZrSiO4 phase forms at 1100  C which is much lower than the former and borosilicate is gone at all investigated temperatures. Therefore, a dense and continuous oxide  layer adhered to SiBCN matrix tightly consisting of fused SiO2 with uniformly distributed ZrSiO4 are observed. The weight gain per unit surface area of the resulting materials as a  function of oxidation  time  is shown  in Fig. 13. As  indicated in Fig. 13(a), a near para-linear weight gain kinetics have been observed. At 1500  C, the oxidation kinetics of SZ3T1800, SZ3T1900 and SZ3T2000 ceramics can be divided  into two parts, parabolic kinetics at  initial 3 h and  linear kinetics thereafter. That  is to say, oxidation behavior of the ceramics deviates from parabolic kinetics and follows a linear law after 3 h of oxidation time. The para-linear kinetics of ZrB2 -SiC ceramics oxidized at 1500  C is reported before [50], which  is matched with our result. At a relatively  low oxida tion temperature of 1100-1300  C, a dense and continuous oxide layer comprising of silica glass and ZrSiO4 covers on oxide surface could prevent the diffusion of oxygen into inner matrix. From the oxide surface morphologies shown  in Fig. 10,  it can be predicted that the oxidation of ceramics is controlled by diffusion control due to dense oxide layer forms at low oxidation temperature. With oxidation temperature reaching or exceeding to 1500  C, phases such as ZrB2 , ZrN, SiC and BNC may be oxidized violently accompanied by the release of gaseous products. This reveals that the growth rate of oxide layer changes from oxygen diffusion through the protective layer to reaction between ZrB2 , ZrN, SiC, BNC with oxygen. However, it seems that the oxidation kinetics of the resulting ceramics, such as SZBT1900 and SZ1T1900 ceramics, are different with high volume content of ZrB2 -ZrN containing ceramics, as  indicated  in Fig. 13(b). SZBT1900 and SZ1T1900  ceramics  show a parabolic kinetics at 1500  C, while SZ2T1900 and SZ3T1900 ceramics reveal a para-linear raw. For higher content of SiC exists in SZBT1900 and SZ1T1900 ceramics,  it can be oxidized  to  form substantial SiO2  \\x0c', 'D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410   4407  Fig. 11.  (a) Surface and  corresponded to (c).  (b) cross-sectional mophologies of SBZT2000 ceramics oxidized at 1500   C   for 5 h,   (c) enlarged   image of dense area   in   (a) and   (d) EDS results  Fig. 12.  (a) Gibbs free energy and (b) reaction enthalpy changes of reactions predicted occurred on oxide surface  in a range of 300-2000 K, phase diagrams of (c) B2 O3 -SiO2 and (d)SiO2 -ZrO2 .  on oxide  scale, which can act as a barrier against  inward oxygen transport. Thus, a dense SiO2 glass layer with a small amount of uniformly distributed ZrSiO4 provides  the positive oxidation behavior with a parabolic increase in oxidation thickness. Furthermore, the weight gain per unit surface area of high volume content of SiC incorporating ceramics is much lower than others. The above analyses are in well agreement with oxide layer morphologies. In summary, at the initial stage of 0-3 h for all the investigated ceramics, namely parabolic kinetics stage,  is controlled by the diffusion of oxygen through the dense oxide  layer consisting of SiO2 glass and uniformly distributed ZrSiO4 . The formation of relatively  low viscosity SiO2 on oxide surface steps from this process, and it will  Fig. 13. Weight variation versus oxidation time for the SiBCN/ZrB2 -ZrN ceramics oxidized at 1500   C in flowing air.  \\x0c', '4408   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  Fig. 14. Surface and cross-sectional mophologies of SZ3T2000 ceramics oxidized for 5 h at different temperatures, inserted enlarge images corresponding to the interface between oxide layer and ceramic matrix: (a), (d) 1100   C; (b), (e) 1300   C; (c), (f) 1500   C.  ×  spread on oxide surface of the ceramics rapidly. Therefore, the partial pressure of oxygen  is  lower than the  inner part of ceramics, which promote the active oxidation reaction (3) happening and SiO dissociation from SiO2 (SiO vaporizing below pO2   10−13 Pa  8.8  at 1500  C) [47,50,51]. At the second oxidation stage above 3 h, the oxide kinetics is depended on composition design of raw materials. Fig. 14 presents the surface and cross-sectional microstructures of SZ3T2000 ceramics by SEM, inserted are the images corresponding to interface between oxide layer and ceramic matrix. From the oxidized exterior of SZ3T2000 ceramics oxidized at 1100  C for 5 h, no visible pores and cracks are  found on oxide  layer and  inter80  face. A thickness of  is obtained and the enlarged  image inserted  indicates  that  the well crystallized phase consists of Si, O and a small amount of Zr  (the  relative EDS  result not shown here). This result shows that the interface between matrix and the outer oxide  layer comprises of mainly SiO2 and a small amount of ZrO2 . With the increase of oxidation temperature, the thickness of oxide  layer enhances rapidly and pores can be seen on oxide surface and  interface between ceramics matrix and outer oxide layer, as  indicated  in Fig. 14(b)-(f).  In Fig. 14(c), the structure of oxide layer can be divided into three layers: (1) SiO2 glass-rich layer with uniformly distributed ZrSiO4 , (2) porous layered structure of  \\u242em   SiO2 -ZrO2 , (3) unreacted layer. However, no SiC-depleted layer are observed, which is a little different with Rezaie’s work [3,7,15]. The formation of SiC-depleted  layer  in ZrB2 -ZrC system depends not only on the surrounding pressure and temperature condition but also on the volume distribution of SiC in ZrB2 matrix [52]. In some systems, such as SiC-ZrB2 and ZrC-SiC ceramics, SiC-depleted layer is not formed due to the volume ratio of SiC/ZrB2 and SiC/ZrC, SiC distribution and the inner oxygen partial pressure [50,53]. The volume ratio of SiC/ZrB2 -ZrN  in this study cannot be defined clearly because ZrN  is  formed by residual Zr and N2 . The oxygen partial  10−13 Pa, considering the pO2 pressure in this region is above 8.8  for the ZrB2 -ZrO2 equilibrium at 1500  C (boundary condition of reaction (3) depends on the pO2 ) [50]. SiC is oxidized to form SiO2 glass rapidly despite the oxidation rate of SiC is much slower than ZrN and ZrB2, showing that the reaction (5) is more dominant than reaction (3).  It can be predicted that the amount of SiO2 phase  in this region would increase by the internal oxygen partial pressure. Fig. 15 shows a SEM image taken from the region slightly below the interface between the matrix and the second oxide layer. Several ceramic samples are prepared to find this interface at varying depth. As indicated, the unreacted areas (white areas) in this image consist of Zr, B, C and N elements which may compose of ZrB2 , ZrN  ×  \\x0c', 'D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410   4409  Fig. 15. Surface morphology of interface between oxide layer and ceramic matrix of SZ3T2000 ceramics oxidized at 1500   C for 5 h and corresponding elements maps. (Several ceramic samples are prepraed to find this interface at varying depth through polishing the oxide layer by hand).  and BNC phases. However, the gray regions reveal that SiO2 is the main phase  in these areas. Combination with the above analyses, we can confirm that the  interface near to the second oxide  layer composes of SiO2 and ZrO2 , while interface close to matrix consists of unreacted area (ZrB2 , ZrN and BNC phases) and a small amount of SiO2. In summary, the oxidation process and formation of oxide layer 1100  C, structure are as  follows. Oxidation at  low temperature  SiC  is oxidized  to  form a SiO2 glass with ZrSiO4 uniformly distributed on surface despite  the oxidation driving  force of SiC  is much lower than others. In this temperature regime, the improvement of oxidation resistance performance aroused by more SiC can be ascribed to the formation of more silica glass and zircon phase ZrSiO4 . In addition, the diffusion energy of oxygen in this oxide scale is higher than pure ZrB2 or ZrN ceramics. However, the morphologies of oxide  layer depend on  the volume ratio of SiC/ZrB2 -ZrN. For  low volume  ratio of SiC/ZrB2 -ZrN containing ceramics, such as SZ2T1900, SZ3T1800, SZ3T1900 and SZ3T2000 ceramics, lots of ZrSiO4 uniformly distribute on oxide surface and  form an  island structure, as indicated in Fig. 10. Due to the low oxidation temperature, the oxidation rate is slight and thickness of oxide layer is thin. At all the oxidation temperatures, borosilicate glass  is destroyed completely due  to  the rapid evaporation of B2O3 and  thermally stable ZrSiO4 forms. Owe to the  formation of a stable SiO2 glass with uniformly distributed ZrSiO4 , the SiBCN/ZrB2 -ZrN ceramics with different volume  ratios of SiC/ZrB2 -ZrN are considered  to show positive oxidation resistance with parabolic mass gain kinetics or para-linear raw, which is similar to the literatures [14,50]. For SZBT1900 and SZ1T1900 ceramics, they exhibit a parabolic kinetics at 1500  C due to the formation of dense and continuous oxide layer which inhibits the oxygen diffusion into inner matrix. Others ceramics, such as SZ2T1900, SZ3T1800, SZ3T1900 and SZ3T2000 ceramics, display a para-linear mass gain raw. Similar to ZrB2 -SiC ceramics  [2,14], a  layered structure  forms during oxidation process of SZ3T2000 ceramics after oxidation test at 1500  C for 5 h in  flowing air. The layered structure consists of three layers: (1) SiO2 glass  layer with uniformly distributed ZrSiO4 (SiO2 -ZrSiO4 ), (2) a layer of SiO2 and ZrO2 (SiO2 -ZrO2 ), (3) unreacted matrix. For high content of ZrB2 -ZrN constituent  in ceramics, SZ3T2000 ceramics suffer a more severe oxidation because of gas evaporation obvious and porous structure of ZrO2 formation which cannot offer an effective anti-oxygen diffusion barrier. We can forecast that high volume ratio of SiC/ZrB2 -ZrN containing ceramics would do better performance on oxidation resistance than others at the same oxidation condition. Therefore, we can prepare SiBCN/ZrB2 -ZrN ceramics combination of  reasonable mechanical properties and excellent oxidation resistance which all depends on our practical needs.  4. Conclusions  Mechanical  activation  assisted  self-propagating  hightemperature  synthesis  (MA-SHS)  plus  reactive  hot  pressing (RHP)  techniques  are  an  effective way  to prepare dense bulk SiBCN/ZrB2 -ZrN  ceramics.  The  disordered  carbon with  amorphous state and mechanical activation should be responsible  for the  ignition of SHS process. As-obtained  ceramics prepared by reactive hot pressing have a similar structure which consists of ZrB2 , ZrN, SiC, BNC and a small amount of zirconium oxide. And ZrN phase comes  from the reaction between residual Zr and N2. Incorporation of high content of UHTCs component perform better on mechanical properties and density, and the flexural strength, fracture toughness, Young’s modulus and Vicker’s hardness of the optimal  result  reach  to  400.0   15.4 MPa,  3.16   0.19 MPa m1/2 , 251.6   5.8 GPa and 9.57   1.24 GPa, respectively. After oxidation, the oxide layer of resulting ceramics composes of SiO2 , ZrSiO4 and ZrO2 with a  layered structure. The  formation of SiO2 glass with uniform distribution of ZrSiO4 on oxide surface acts as an effective anti-oxygen  diffusion  barrier,  hence  performs  an  outstanding oxidation  resistance. The oxidation behavior of SiBCN/ZrB2 -ZrN ceramics in structure changes is similar, but the oxidation kinetics  ±  ±  ±  ±  \\x0c', '4410   D. Li et al. / Journal of the European Ceramic Society 35 (2015) 4399-4410  is different. The higher content of SiC containing ceramics exhibit a parabolic kinetics while others show a para-linear kinetics at 1500  C.  In flowing air at 1500  C,  ceramics with high  content of UHTCs  component  show passive oxidation with para-linear kinetics duo to the porous oxide layer structure resulting from SiC, ZrB2, ZrN and BNC oxidized beneath the outermost layer and gases evaporation. 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Telle, Thermal stability of zircon (ZrSiO4 ), J. Eur. Ceram. Soc. 28 (2008) 2199-2211. [50] Y.H. Seong, S.J. Lee, D.K. Kim, TEM study of the high-temperature oxidation behavior of hot-pressed ZrB2 -SiC composites, J. Eur. Ceram. Soc. 96 (2013) 1570-1576. [51] C. Tian, D.Y. Gao Zhang, C. Xu, Y. Song, X. Shi, Oxidation behavior of zirconium diboride-silicon carbide ceramic composites under low oxygen partial pressure, Corros. Sci. 53 (2011) 3742-3746. [52] X.H. Zhang, P. Hu, J.C. Han, Structure evolution of ZrB2 -SiC during the oxidation in air, J. Mater. Res. 23 (2008) 1961-1972. [53] L.Y. Zhao, D.C. Jia, X.M. Duan, Z.H. Yang, Y. Zhou, Oxidation of ZrC-30 vol% SiC composite in air from low to ultrahigh temperature, J. Eur. Ceram. Soc. 32 (2012) 947-954.  \\x0c']"
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  "_id": 213,
  "PDF": "Pressureless sintering of (ZrB2–SiC–B4C) composites with (Y2O3+Al2O3) additions.pdf",
  "Text": "['Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  Contents lists available at ScienceDirect  Int. Journal of Refractory Metals and Hard Materials  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / I J R M H M  Pressureless sintering of (ZrB2-SiC-B4C) composites with (Y2O3 + Al2O3) additions  R.V. Krishnarao ⁎, Zafir Alam, D.K. Das, V.V. Bhanu Prasad, G. Madhusudan Reddy  Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad 500058 India  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 1 April 2015 Received in revised form 13 May 2015 Accepted 18 May 2015 Available online 21 May 2015  Keywords:  ZrB2 Composites Pressureless sintering Oxidation Thermal shock Joining  1. Introduction  Pressureless sintering of (ZrB2-SiC-B4C) composites with (Y2O3 + Al2O3) additions has been studied. The vol.% of SiC, B4C, and (Y2O3 + Al2O3) was varied from (26, 24, and 16) to (5, 4, and 5) respectively. Hardness of composites has been found to decrease with a decrease in B4C content. Flexural strength has been found to increase with decreasing (Y2O3 + Al2O3) content. No considerable decrease in flexural strength for composite with high vol.% of SiC, B4C, and (Y2O3 + Al2O3) has been observed up to a temperature of 1500 °C. By optimizing the composition a composite possessing a density of 4.64 g cm−3, flexural strength of 213 MPa, and Vickers hardness of 17.3 GPa has been sintered. Conical shapes could be made by manual shaping and sintering. A complex yttria-aluminasilicate (YAG) layer was found to protect the composite from oxidation at high temperature of 1700 °C. The composites exhibited good dimensional stability and thermal shock resistance at 2200 °C in oxy-acetylene flame and at 2700 °C in plasma flame. Formation of yttria stabilized zirconia embedded in the matrix of YAG has been identified on the flame exposed surfaces. The composite could be joined to itself by gas tungsten arc welding (GTAW) with a filler material containing (ZrB2-SiC-B4C-YAG). The shear strength of the weld was about 50% of the flexural strength of the parent composite.  © 2015 Elsevier Ltd. All rights reserved.  Borides, carbides and nitrides of transition elements such as hafnium, zirconium, tantalum and titanium possess a unique combination of ultra-high melting points (N 2500 °C), high strength, high hardness and high electrical and thermal conductivities. These ultra-high temperature ceramics (UHTC) are considered for extreme environments that require oxidation resistance above 2000 °C [1-3]. Among them ZrB2 with the lowest theoretical density is an attractive material for extreme thermal and chemical environments associated with hypersonic flight, rocket propulsion, and atmospheric re-entry. ZrB2 is being considered for high speed air craft leading edges, as well as for structural parts in high temperature environments [4-6]. Because of the difficulty in obtaining full density due to its covalent bonding, and low self diffusion co-efficient, ZrB2 based composites have been densified by hot-pressing (HP) [7-9]. HP processing is limited to the formation of simple geometries and moderate sizes. Further diamond machining to fabricate complex shapes is an expensive and time consuming. Low cost near net-shape processing of ceramic parts with complex geometries is possible with pressureless sintering (PS) as it will minimize machining.  ⁎  Corresponding author. E-mail address: rvkr4534@yahoo.com (R.V. Krishnarao).  http://dx.doi.org/10.1016/j.ijrmhm.2015.05.013 0263-4368/© 2015 Elsevier Ltd. All rights reserved.  A variety of sintering additives to enhance the densification of ZrB2 ceramics can be classified into: liquid formers, solid solution formers, and reactive agents. Liquid phase formers include refractory metals such as Ni, Fe, Co, and Mo [10-12], as well as disilicides of Mo [13,14] and Zr [15]. Other refractory metals such as Nb and Cr [16] have also been used to enhance PS of ZrB2. ZrB2 can form solid solutions with Ti, Hf, Nb, Ta, and Mo borides over a wide range of compositions and improve the densification of ZrB2 [17]. Reactive agents react with oxide impurities (ZrO2, B2O3) present in the surface of starting particles which inhibit densification. The main reactive agents used widely include B4C, C, and WC [18-20]. The major problem with PS of ZrB2 is exaggerated grain growth and low mechanical properties. Using combination of B4C and C, ZrB2 could be sintered to almost full density at 1900 °C in 120 min [21]. The addition of appropriate amounts of SiC to ZrB2 strongly improves its oxidation resistance at high temperature [22]. SiC is always added to ZrB2 to improve the densification, oxidation resistance, fracture toughness and mechanical strength [23-26]. The addition of rare earth oxide yttria (Y2O3) has been shown to react with oxide impurities (ZrO2, B2O3, SiO2) in ZrB2-SiC composites and improve densification and mechanical properties [27]. Similarly Al2O3 + Y2O3 have been used to reduce the sintering temperature and to increase densification rate of ZrB2-SiC composites [28]. The formation of low melting grain boundary phases from sintering aids often deteriorates the high temperature properties. Crystallization annealing  \\x0c', '56  R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  can result in YAG phase from grain boundary and enhance high temperature properties of ZrB2-SiC ceramics [29]. YAG (Al2O3 + Y2O3)/ZrB2 composites with high density and oxidation resistance could be sintered at relatively low temperature of 1600 °C [30]. The difficulty in fabricating large size or complex shapes limits the application of ZrB2-SiC based composites. Joining them by fusion welding without or with preheating, controlled cooling under protective gas shield lead to thermal shock failure or porosity at the weld interface [31,32]. There exists the necessity to develop a suitable filler composite material possessing oxidation resistance and thermal shock resistance to join ZrB2-SiC based composites. In the present study (ZrB2-SiC-B4C) composites were pressureless sintered at relatively low temperature between 1550 and 1680 °C using (Al2O3 + Y2O3) additions to produce an oxidation resistant, thermal shock resistant, and weld able composites for high temperature applications.  2. Experimental  2.1. Materials  Pressureless sintering of ZrB2-SiC-B4C composites has been carried out using two different types of powders of ZrB2. The self-made ZrB2 (SM) powder used in this work was synthesized by reacting ZrO2 with B4C according to reaction:  7ZrO2 þ 5B4 C→7ZrB2 þ 5CO þ 3B2O3 :  ð1Þ  To obtain a single phase ZrB2 without impurities like un-reacted ZrO2 , B4C , and free C , the excess o f B4C was taken in wt . ratio o f ZrO2/B4C = 2.5 where the stoichiometric wt. ratio was ~ 3.0. The details of synthesis were reported elsewhere [33]. ZrO2 powder supplied by Nuclear Fuel Complex , Hyderabad , India and B4C powder supplied by China Abrasives, Zing Zhou, China were used (Table 1). The average particle size of the as synthesized ZrB2 was b 1 μm. The typ ica l chem ica l ana lys is o f the ZrB2 is given in Tab le 2 . Another source of ZrB2 was the Grade A powder with a particle size (d50 ~ 3.0- 5.0 μm) supplied by HC Starck, Germany. SiC powder with a particle size (d50 ~ 0.8 μm) was supplied by H.C. Starck, Germany. Al2O3 with a super fine size (d50 ~ 0.7 μm) obtained from Alcan and sub-micron sized Y2O3 were used. Hot pressing at 2000 °C of a mixture of powders containing equal volumes of ZrB2, SiC, and B4C has been reported to yield a composite with super hardness of 28.9 GPa (~ 43 GPa), low bulk density of 4.04 g cm− 1, and with fracture strength of 603 +/− 155 MPa [20]. In this work we have added initially about 8 wt.% each of Y2O3 and Al2O3 to enable low temperature PS of the similar ZrB2, SiC, and B4C composition. After studying the mechanical, oxidation and high temperature properties of initial composite four more composites by decreasing the B4C and YAG contents have been prepared. The initial vol.% of different powders (ZrB2, SiC, B4C, Y2O3 and Al2O3) and designation of respective compositions are given in Table 3.  Table 1 Purity of reactants used for synthesis of ZrB2.  Chemical Impurity  ZrO2  B4C  C.Black  H f b 50 , A l b 10 , B b 0 .5 , Ca b 125 , Co b 10 , Cr b 25 , Cu b 25 , Fe b 100 , Mg b 25 , Mn b 5 , Mo b 25 , N i b 15 , Pb b 25 , T i b 50 , Sn b 10 , V b 10 ppm B4C 92.65%, total B 76.25%, compound B 75.0%, free B 1.25%, Fe2O3 0.45%, B2O3 0.12%, Total C 17.65%, compound C 17.65%, free C 2.56% Grade N774  Table 2 Analysis of self made ZrB2 and HC Starck Grade A ZrB2.  Self-made ZrB2  Carbon Oxygen Boron Zirconium  Wt.%  0.63 0.23 17.96 81.18  HC Starck ZrB2  Carbon Oxygen Boron Zirconium  Wt.%  0.15 0.50 19.00 80.35  2.2. Pressureless sintering  The dry mixing of powders was done for 24 h in a polythene bottle with alumina balls. Green compacts of 60 mm diameter were made using PVA binder in water solution and uni-axial compaction with a load of 9-10 Tons. Few SM1 samples were pressed manually in steel die by mild hammering of the ram. Cone shaped green compacts of SM1 were also made using mold made out of filter paper or by manually forming without using any mold. The pressureless sintering was carried out in a graphite resistance heating furnace (ASTRO, USA, Model 10003060-FP20). The furnace was evacuated to a moderate vacuum (5 × 10− 2 Torr) and back filled with argon up to a pressure of 1 atm. after reaching a temperature of 1020 °C to facilitate debinding in a vacuum. Temperature was maintained with a Model 939A3 Honeywell radiation pyrometer. Heating rate employed was ≈ 15 °C min− 1. The composites were pressureless sintered at 1680 °C in argon atm. for 1.0 h. Holding for a minimum time of 0.5 h at 1550 +/− 25 °C was employed to avoid total melting of the sample when directly heated to above 1600 °C. Heating rapidly to above 1600 °C can lead to melting and sticking of the sample to graphite crucible due to the formation of large quantity of liquid phase. Sintered compacts of composite of size 10 mm diameter × 10 mm height and 30 mm diameter × 10 mm height were also made by pressureless sintering. A 30 mm diameter and 10 mm thick monolithic ZrB2 compact with Grade A powder supplied by HC Starck, Germany was also sintered to study the oxidation upon exposure to plasma flame of 2700 °C.  2.3. Preparation of filler material for arc welding  Though, arc welding of ZrB2 is possible research on joining by GTAW is very limited. Earlier efforts on fusion welding of ZrB2-20 vol.% SiC [31] and ZrB2-20 vol.% ZrC composites [32] by pre-heating and controlled cooling under protective atmosphere lead to thermal shock failure or Suitable filler porosity at the weld interface. composite material possessing oxidation resistance and thermal shock resistance may be useful to join ZrB2-SiC based composites. Since the filler material undergoes melting and flowing into the weld gap it should have more flow ability and oxidation resistance compared to parent material. After studying the oxidation behavior SM1 composite was selected for this purpose. The dry mixing of powders was done for 24 h in a polythene bottle with alumina balls. Thick paste was made using PVA binder in water solution. The paste was filled into a syringe used to inject medicine and extruded without using needle to get a ~ 3 mm diameter and 7.5-10 cm long rods. After initial natural drying the filler rods were dried in an oven at 110 °C for 1 h. The pressureless sintering was carried out at 1650 °C in argon atm. for 1 h.  2.4. Gas tungsten arc welding  The bar samples of size 4 × 5 × 50 mm long were used for GTAW of SM4 composites to themselves. The samples were ultrasonically cleaned in acetone and dried in air before fusion welding. The weld coupons were laid on a steel platform. A welding current of 90-120 A and manual welding speed of 3 mm/min were used without any pre-heating of weld coupons. The butt weld gap was ~ 1 mm. After welding the argon flow was continued till the joint was cooled to a temperature  \\x0c', 'R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  57  Table 3 Designation and mechanical properties of different composites with self-made (SM) ZrB2 and HC Starck (HC) ZrB2.  Designation  Vol.%  Calculated density g cm− 3  Measured density g cm− 3  Open porosity %  VHN GPa @ 200 g  Flexural strength MPa  ZrB2  SiC  B4C  Y2O3  Al2O3  SM1/HC1 SM2/HC2 SM3/HC3 SM4/HC4 SM5/HC5 HCP  34 50 59 65 86 100  26 29 26 20 5 0  24 10 6 8 4 0  7 5 4 3 2 0  9 6 5 4 3 0  4.20 4.72 4.96 5.10 5.70 6.10  3.70/- 3.80/4.22 4.36/4.39 4.64/4.59 -/4.73 4.40  0.90/- 0.58/0.87 0.78/1.90 0.72/1.90 -/4.19 16.00  25.39/- 20.32/15.89 12.19/11.82 17.30/14.83 -/10.06  -  112/- 143/120 165/113 213/140 -/098  -  Table 4 Flexural strength of SM1 composites at different temperatures.  Sample  Flexural strength MPa  SM1 machine pressed in die SM1 manually pressed in die  RT  112 111  1200 °C  1500 °C  108 113  85 111  less than 800 °C. Similarly the arc welding was performed on the opposite side of the weld. GTAW machine of ESAB make, Model TIG 300A, Kolkata, India was used.  2.5. Characterization  The bulk density was calculated using a water displacement method. The joints were cut perpendicular to the welding direction using a diamond cutting wheel or CNC wire cut EDM. The cut welds were mounted in epoxy and polished to mirror finish using a fine diamond (0.25 μm) abrasive. Similarly three point bend specimens of size 4 × 5 × 50 mm, and sample coupons of size 10 × 10 × 2 mm for oxidation studies were prepared. Instron of model No: 8801 was used to measure the flexural strength with a span of sample of 40 mm and cross-head speed of 0.5 mm/min. The SM1 samples were tested at room temperature, and high temperature of 1200 °C and 1500 °C. Isothermal oxidation was carried out for 15, 30, 45, and 60 min at different temperatures between 1100 °C and 1700 °C. The oxidation in air was studied using a raising hearth furnace of Naskar & Co., Kolkata, India. The 30 mm diameter and 10 mm height samples were subjected to flame exposure to oxy-acetylene flame at 2200 °C, and plasma flame at 2700 °C. The temperature of flame and sample was measured with a precision optical pyrometer supplied by Pyrometer Instrument Co., Inc., USA. In the case of oxy-acetylene flame the samples were exposed for 1-1.5 min before the withdrawal of the flame to allow the sample to  cool naturally for 1.0 min. The samples were exposed for 10-20 times. In the case of plasma flame the samples were exposed continuously for a minimum of 2 min to a maximum of 10 min. The cross sections of polished samples were analyzed for microstructure using an optical microscope and scanning electron microscope (SEM of FEI Quanta 400, Netherlands). The Vickers micro-hardness was measured with DM H-2, Matsuzawa Seiki, Japan, using a load of 200 g for a dwell time of 15 s. A Philips X-ray diffractometer, Model PW3710, with Cu Kα radiation through a Ni filter was used to identify different phases in composite. Shear test specimens from the weld interface region were extracted as per ASTM A 264 standard. The shear strength of the weld was measured using a stainless steel fixture in Walte-BaiAg, HTV-1200, universal testing machine. A cross-head speed of 0.1 mm min− 1 was used. The shear strength was calculated from the maximum load value divided by the overlap area.  3. Results and discussion  3.1. Mechanical properties  The measured bulk densities varied from 3.70 g cm− 3 to 4.73 g cm− 3, for SM and HC composites (Table 3). The large difference in densities from calculated values to measured values could be due to the loss of oxide vapors of all constituents of composites. Considerable weight loss (N 10%) for different composites was observed after sintering. About (~ 5%) weight loss was observed for the monolithic ZrB2 sample sintered under similar conditions. During liquid phase sintering under free flow of argon, loss of gaseous species is possible. In the presence of many phases (such as ZrB2, SiC, and B4C), additives (Y2O3, Al2O3) and surface oxides (SiO2, ZrO2, B2O3) the reaction among them during sintering and cooling is unknown and the densification mechanisms may be controlled by dissolution-precipitation or  a   Y  * ZrB  2   SiC     B C    4   YAG   *              *                 *  s  z  B   15µm   Cu Kαα 2θ deg  Fig. 1. Typical XRD pattern of SM1 sintered at 1680 °C.  Fig. 2. Typical BSE image of SM1 sintered at 1680 °C. Bright phase — ZrB2, gray phase — SiC, dark phase — B4C, and white crystalline phase — YAG.  \\x0c', '58  R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  Sintered           1500 C              1600 C            1700oC     o  o                      900 s                      3600 s                                               Sticking   Fig. 3. Appearance of SM1 samples: as sintered and after oxidation for 15 min and 60 min at different temperatures.  evaporation-condensation [27]. A sintering temperature of 1600 °C is considered as a critical temperature to form molten YAG and to achieve a full density of YAG/ZrB2 composite by liquid phase sintering [30]. In the present study rapid heating to temperatures above 1600 °C has led to total melting of the sample and reaction with graphite crucible. Since the true densities of the sintered composites are not known the comparison of relative densities based on the measurement of bulk densities is only an indication of the trend of increasing or decreasing [34]. Bulk densities are in the range of 88-90% of calculated densities. Irrespective of the material composition the open porosity is 0.9% to 2.0%. The highest open porosity of 4.19% was recorded for HC5 sample as the SiC content was the lowest (5 vol.%) in this sample. The monolithic ZrB2 sample sintered under similar conditions has attained about 72.13% of bulk density and 16.00% of open porosity. Vickers hardness number (VHN) decreased with B4C content for both SM and HC samples. For similar compositions, the VHN for HC samples was lower than that of SM samples (Table 3). This could be due to difference in composition (Table 2), synthesis route, and particle size of ZrB2. As the SM ZrB2 is synthesized through B4C reduction of ZrO2 the high residual carbon and low oxygen content may be contributing to the higher VHN. The lowest flexural strength 112 MPa was obtained for SM1. No considerable decrease in flexural strength has been noticed at 1200 °C and 1500 °C for SM1 (Table 4). The manually pressed SM1 samples have also retained the room temperature flexural strength of 111 MPa at 1200 °C and 1500 °C (113 MPa and 111 MPa respectively). The flexural strength of ZrB2-SiC composites at high temperature is strongly influenced by the grain size and SiC content. Decreasing the particle size is beneficial to the room temperature strength but detrimental to the ultra-high temperature strength. Grain boundary sliding and cavitations are the controlling flexural deformation mechanism in ZrB2-30 vol.% SiC at ultra-high temperature. The flexural strength of ZrB2-15°vol.% SiC with large starting particle sizes (5 μm and 2 μm respectively) at 1800 °C was 217 MPa against its room temperature strength of 500 MPa [35]. In the present case the retention of strength up to a temperature of 1500 °C may be attributed to a large sintered grain size. The typical XRD pattern of the SM1 (ZrB2-SiC-B4C) sintered with YAG (Al2O3 + Y2O3) is shown in Fig. 1. The major phase is ZrB2 and other phases SiC, B4C and YAG were identified. In the backscattered electron image (Fig. 2) the bright phase is ZrB2, gray phase is SiC, dark phase is B4C and white crystalline particles are identified as YAG phase. The retention of strength at high temperatures of 1200 °C and 1500 °C could be due to the large grain size and crystallization of YAG phase. In (ZrB2-SiC-B4C) composites sintered by spark plasma sintering (SPS) without additions of (Al2O3 + Y2O3) the increase in B4C content  from 1 wt.% to 5 wt.% has been reported to decrease in density and increase in porosity which can be a source of crack initiation. In spite of the decrease in room temperature strength the strength at high temperature was shown to increase due to filling of pores with glass formed by the reaction of ZrO2 on ZrB2 with B4C [36]. Further to increase the room temperature flexural strength the vol.% of (Al2O3 + Y2O3) has been decreased. With a decrease in YAG (Al2O3 + Y2O3) content the flexural strength increased to 213 MPa and 140 MPa for SM4 and HC4 respectively. The flexura l strength of 98 MPa was the lowest for HC5 in which the vol.% of all additives inc lud ing S iC was very low . Among the number o f reports on PS o f ZrB2-20 vol.% SiC, few reports [29,37-39] that utilized sintering additives (Y2O3-Al2O3/B4C/C/phenol resin) can be compared with the present compos ites . The reported flexura l strength va lues var ied from 202 MPa to 315 MPa, which is in agreement with the present va lue of 213 MPa obta ined for PS compos ites of (ZrB2-S iC-B4C- YAG).  3.2. Oxidation up to 1700 °C  After oxidation up to a temperature of 1300 °C there is no change in the appearance of the sample (Fig. 3). After oxidation at 1500 °C for 0.25 h a slight change in appearance due to the formation of small white particles is observed. With an increase in oxidation temperature  Fig. 4. Plot of weight change with time of SM1 samples oxidized at different temperatures.                 \\x0c', 'R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  59  and time the samples appeared white with glossy texture. Oxidation kinetics measurements are generally based on weight change or oxide scale thickness changes with time upon exposure to oxidizing atmosphere at different temperatures. Both methods have limitations due to simultaneous oxidation (formation of ZrO2, B2O3, and SiO2) and vaporization (loss of CO, B2O3, and SiO2) processes are taking place. To overcome this limitation measurement of total oxygen consumption per unit area of sample or evaluation of micro-structural changes can be utilized. The weight change after oxidation for 1 h at 1100 °C, 1300 °C, and 1500 °C was very negligible i.e. less than 5 mg cm− 2 (Fig. 4). At 1600 °C after 0.5 h and at 1700 °C after 0.25 h the samples were sticking to Al2O3 crucible. The rapid weight loss recorded was attributed to the reaction of oxide glass with Al2O3 crucible. The formation of YAG layer on the oxidized samples was identified through energy dispersive spectroscope (EDS) analysis. ZrB2 has been investigated for its oxidation performance since 1960s [40,41]. During this period the reported findings were reviewed by  many researchers [42-44]. Upon heating in air the monolithic ZrB2, oxidizes to form skeleton of ZrO2 and B2O3 liquid at temperatures as low as 450 °C.  2ZrB2  sð  Þ þ 5O2 gð Þ→2ZrO2  sð  Þ þ 2B2O3  lð Þ:  ð2Þ  The liquid B2O3 wets the ZrO2 grains until it volatilizes at temperatures above 1100 °C. Below this temperature the diffusion of oxygen through the liquid B2O3 that surrounds the ZrO2 grains is the rate limiting step . Between 1100 °C and 1400 °C , the mass gain due to the formation of ZrO2 and B2O3 and mass loss due to vaporization of B2O3 occur.  B2O3  lð Þ↔B2O3 gð  Þ  ð3Þ  When S iC is added i t fac i l i ta tes l iqu id pha se s in te r ing v ia the fo rma t ion o f a bo ros i l ic a te l iqu id . Du r ing ox ida t ion a t e leva ted temperatures , the s i l icon and boron in (ZrB2-S iC) are ox id ized to  Fig. 5. BSE images of surfaces of SM1 samples oxidized at (a) 1100 °C/60 min, (b) 1300 °C/60 min, (c) 1500 °C/15 min, precipitation of ZrO2 from glass, (d) 1600 °C/15 min, formation of large blisters on the oxide layer, (e) 1600 °C/15 min, debris in collapsed blister on the oxide layer, (f) 1600 °C/60 min, formation of small cavities in oxide layer, (g) and (h) EDS of oxide surface in (a) and (b) respectively, and (i) typical EDS of white globules in (c) and (f).  \\x0c', '60  R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  form a protective borosilicate glass layer according to Eqs. (3) and (4) .  SiC sð  Þ þ 3=2O2→SiO2  lð Þ þ CO gð Þ  ð4Þ  The diffusion coefficient of oxygen through borosilicate is lower than those for diffusion in B2O3 and ZrO2. It has higher viscosity, higher boiling point, and efficient oxidation protection. Above 1400 °C the volatilization of B2O3 is higher than the production of B2O3. The weight increase due to the ZrO2 formation is greater than the weight reduction due to ZrB2 consumption and to the evaporation of B2O3. Borosilicate glass dissolves ZrO2 and forms BSZ channels that connect the interface between the oxide layer and the unoxidized substrate. This BSZ glassy phase can act as a carrier of gas species (O2, CO, CO2, and SiO) that can diffuse in it for the oxidation of the ZrB2-SiC material bulk [45,46]. Due to progressive evaporation of B2O3 the microstructure of the oxide scale changes with the temperature. The evaporation of B2O3 increases the viscosity of the glassy phase and decreases the solubility of ZrO2, then secondary ZrO2 precipitates on the external part of the scale. The oxide scale is considered as a multi-phase layer changing its composition, physical properties (like viscosity), and relative amount of phases with the temperature. In the present investigation the micro-structural changes occurred on the sample surfaces upon oxidation in air at different temperatures are shown in backscattered electron (BSE) images in Fig. 5. After oxidation at 1100 °C for 1 h the electron dispersive spectroscopy (EDS) analysis revealed the presence of Y, Al, O, and with major element of Si (Fig. 5(a) and (g)). After oxidation at and above 1300 °C the EDS showed the presence of Si, Al, O, Y, and Zr (Fig. 5(b) and (h)). This confirms that in SM1 the borosilicate glass formed at low temperatures dissolves the YAG and primary ZrO2 to form a complex yttria alumina silicate glass containing ZrO2. As the oxidation temperature increased to 1500 °C the evaporation of B2O3 causes the precipitation of ZrO2 from the glass layer (Fig. 5(c)). The minute B2O3 gas bubbles collectively form blisters on the glossy layer and collapse when the pressure inside the blister exceeds the ambient pressure. The appearance and collapse of blisters can  be observed in Fig. 5(d) and (e). Even after rupturing the blister the glass layer is intact and contained the debris of the collapsed blister. As the evaporation of B2O3 continues the viscosity of glass layer increases with temperature. The diameter of the blisters that can form in viscous layer will decrease with an increase in depth. After rupturing of the blister the opening on the surface will remain irrespective of its depth (Fig. 5(f)). The small particles precipitated in the glass layer are identified as ZrO2 (Fig. 5(i)). Further, the evaporation of B2O3 and precipitation of ZrO2 cause an increase in the viscosity of the glass layer and rapidly changes in its physical properties. The volumetric shrinkage causes the cracking of the glassy layer and coalescence of ZrO2 particles (Fig. 6(a) and (b)). The formation of B2O3 gas bubbles and subsequent blisters on the oxide glass surface can be noticed in the cross section of the oxide layer in Fig. 6(c). The agglomeration of ZrO2 particles causes the formation of ZrO2 layer within the glass layer (Fig. 6(d)). The YAG glossy layer together with ZrO2 can protect the composite from further oxidation at higher temperatures for longer durations. A passive to active transition of SiC oxidation (reaction (5)) occurs between 1600 °C and 1700 °C in air under an atmospheric pressure [47,48]. The SiC depletion layer forms below an intermediate layer containing a BSZ glassy phase and ZrO2 particles (Fig. 6(d)).  SiC sð  Þ þ O2 gð  Þ→SiO gð  Þ þ CO gð Þ  ð5Þ  But recent study on the oxidation of ZrB2-SiC compounds with high SiC contents, specifically over the 50 vol.% [49], the formation of a SiCdepleted layer above the ZrB2-SiC bulk composite was not observed. Compositional and micro-structural factors are expected to control the formation of SiC depletion layer in the oxide scale.  3.3. High temperature flame exposure  Further the (ZrB2-SiC-B4C-YAG) composites are exposed to oxyacetylene flame at 2200 °C (Fig. 7(a)) and plasma flame at 2700 °C (Fig. 7(c)). The flame temperature and the sample temperature  (a)   (c)   (b)   (d)    15 µm     50 µm    YAG glossy  layer   YAG + B2O3 bubbles and  ZrO2  SiC depleted  layer   Un-affected    50 µm     125 µm    YAG +  ZrO2  SiC depleted  layer   Un-affected   Fig. 6. BSE images of surfaces of SM1 samples oxidized at (a) 1600 °C (b) 1700 °C SEM images of cross section of oxide layer of SM1 samples oxidized at (c) 1600 °C, and (d) 1700 °C.  \\x0c', 'R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  61  Fig. 7. SM1 samples being exposed to (a) oxyacetylene flame and (c) plasma flame, (b) immediately after withdrawing the oxy-acetylene flame, and (d) cone of SM4 after exposing for 2 min to plasma flame.  immediately after withdrawing the flame (Fig. 7(b)) have been measured. All samples exhibited dimensional stability and thermal shock resistance. The manually made conical sample of SM4 also retained its shape and dimension after exposure for 2 min to plasma flame at 2700 °C. Its weight change was + 8.65 mg cm− 2 (Fig. 7(d)). Since oxidation (formation of ZrO2, B2O3, and SiO2) and vaporization (loss of CO, B2O3, and SiO2) processes are taking place simultaneously the weight gain measurement does not offer good support to explain any oxidation mechanism. But preliminary comparison of oxidation behavior can be made. The pure ZrB2 sample after exposing to plasma flame of 2700 °C for 600 s recorded a weight gain of 7.29 mg cm− 2 (Fig. 8(a)). But SM1 sample recorded a weight loss of 27.87 mg cm− 2 (Fig. 8(b)). In pure ZrB2 the weight loss due to the evaporation of B2O3 is less than that of weight gain due to the formation of ZrO2. In SM1 the sum of losses of CO, B2O3, and SiO2 at 2700 °C is higher than the weight gain due to the formation of ZrO2. After exposure to oxy-acetylene flame the weight change of + 4.0 mg cm− 2, for SM1 is the lowest among all samples tested. SM1 sample retained its shape and dimension even after 20 times exposure to oxy-acetylene flame at 2200 °C (Fig. 8(c)). With a decrease in vol.% of B4C and YAG the weight gain for SM4 has increased to + 11.67 mg cm− 2 after 10 times exposure to oxy-acetylene flame at 2200 °C (Fig. 8(e)). With similar composition HC4 sample gained a weight of 44.5 mg cm− 2 after 10 times of exposure (Fig. 8(d)). Similar difference in Vickers hardness is observed for SM and HC samples (Table 3). This is attributed to the difference in composition (Table 2),  synthesis route, and particle size of ZrB2. As the SM ZrB2 is synthesized through B4C reduction of ZrO2 the high residual carbon and low oxygen content may be contributing to the higher oxidation resistance. The appearance of samples in Fig. 8(d and e), also indicates that the HC4 sample was more oxidized than that of SM4. The XRD analysis of SM1 sample exposed to plasma flame or oxyacetylene flame (Fig. 9) revealed the formation of yttria stabilized zirconia layer on the surface. Passive oxidation protection is expected by the continuous and compact ZrO2 layer which prevents direct exposure of the ZrB2-SiC composite to air. ZrO2 cannot adhere to ZrB2 at high temperature and causes cracking. Spalling tends to occur due to the weak bonding results from the mismatch of coefficient of thermal expansion between the oxide scale and unaltered ZrB2 matrix [50]. In the present composite the YSZ precipitates and remain embedded in complex YAG layer that adheres to parent composite and protects it from further oxidation. Therefore very low weight change values are recorded in the present work. A number of readings were taken from optical photo-micrographs (Fig. 8) and from SEM images (Fig. 10) to measure the average thickness of oxide layers grown after flame exposure for 60 s. Porous and thickest (23.5 μm) oxide scale was observed for HC pure ZrB2 after exposure to plasma flame (Fig. 10(a)). The thickness of the oxide layer was not changed with the change in flame from oxy-acetylene (9.28 μm) to plasma (9.25 μm) for SM1 (Fig. 10(b) and (c)). Their oxide scale thickness was about 40% of that for pure ZrB2. With a decrease in vol.% of (Y2O3 + Al2O3), the thickness of oxide layer was not much decreased  Wt. Change in mg cm-2                   Plasma / 600 seconds Oxy acet. / 1200 seconds Oxy acet. /  600 seconds          +7.29                       -27.87                         +4.0                                +44.5                           +11.67   HC ZrB2                         SM1                         SM1                               HC4                            SM4 selacsedixofossenkciht  Fig. 8. The appearance and optical photo-micrographs of cross section of oxide layers of different samples.  \\x0c', '62  R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65        *  * YSZ                     *              *                                       *   Cu Kαα 2θ deg   Fig. 9. Typical XRD pattern of SM1 surface exposed to plasma flame.  for HC4 sample (Fig.10(d)) whereas a considerable decrease was observed for SM4 (Fig. 10(e)). The oxide scale thickness of SM4 was nearly 1/4 of that for SM1. But the corresponding weight gain for SM4 was higher than that for SM1. The thickness of the glassy oxide layer that can form on particular composite alone cannot control the weight change or oxidation of bulk. The composition of the glass layer on the composite is also important. Additives can improve the oxidation resistance of the ZrB2 composites by (a) increasing viscosity of the borosilicate glass, (b) inhibiting polymorphic transformations of the ZrO2, and (c) modifying the microstructure of the ZrO2 scale. The benefit of modification of the glass composition is limited. During exposure to ultrahigh temperature environments however viscous the liquid phase is, the materials still need to resist the shear forces. ZrB2 based composites  that produce self-generating refractory oxidation barriers or dense ZrO2 scales show the greatest promise. Rupturing due to spalling of the oxide scale in a ZrB2-SiC ceramic after several thermal cycles from room temperature to 1700 °C has been reported [51]. The cracking and spalling of the oxide scale were attributed to a phase transformation in the ZrO2, which takes place during thermal cycling. The oxides also have higher coefficients of thermal expansion (CTEs) and lower thermal conductivities than the underlying diboride material. This, coupled with the phase transformation, leads to cracking of the oxide scale, allowing the oxidation of the underlying bulk. Even after exposing for 10 to 20 times no such spalling or cracking of oxide layer was observed for oxy-acetylene flame exposed samples (Fig. 8(c, d, and e)). In the present composites the formation of YSZ (bright phase in Fig. 10(c)) embedded in YAG (gray phase in Fig. 10(c)) was identified through EDS analysis (Fig. 11(b)). The pure ZrO2 formed on pure ZrB2, undergoes transformation from monoclinic to tetragonal at 1170 °C on heating and at 950 °C on cooling. This leads to the cracking of the oxide scale, allowing oxidation of the underlying ZrB2. The difference in EDS of ZrO2 formed on pure ZrB2 (Fig. 11(a)) and EDS of ZrO2 formed on (ZrB2-SiC-B4C) composites sintered with (Y2O3 + Al2O3) additions can be noticed (Fig. 11(a)).  3.4. Gas tungsten arc welding  Initially SM4 composite bars were tried to fusion weld to themselves by GTAW. The formation of cracks and pores on either side of the joint interface in a heat affected zone was observed. During welding oxidation can induce porosity in the melt fusion zone. Some porosity at boundary of parent material and fusion zone is expected due to volume change upon solidification of melt pool. Thermal shock behavior of a material largely depends upon the number of mechanical and thermo-physical characteristics like elastic    Average thickness of oxide layer / 60 seconds in parenthesis         HCP ZrB2  (23.5 µm)              SM1    (9.25 um)   plasma flame exposed  oxy acetylene flame  exposed              SM1    (9.28 µm)                    HC4     ( 9.05 µm)                    SM4    (2.18  µm)  Fig. 10. BSE images of cross section of oxide layer of different samples after exposure to flame.  \\x0c', \"R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  63  material can also lead to the escape of gaseous species like SiO, CO, and B2O3 and induce porosity in the fusion zone. To avoid cracks, and pores during welding a suitable filler material is required. Filler form a liquid pool and fill the gap between the surfaces to be joined. It is similar to metal casting into a mold. By properly controlling the welding speed the flow of the filler liquid into a weld gap can be controlled to avoid cracks and pores that could form due to shrinkage during the solidification of molten filler. The filler material should have more flow ability and oxidation resistance compared to parent material. It should flow into the weld gap easily. Sintered SM4 compacts of 10 mm diameter × 10 mm height and 30 mm diameter × 10 mm height were joined by GTA melting the filler rods of SM1 composition. The joint interface between the parent material and filler material was very clean and coherent Fig. 12(a). The Vickers hardness across the interface was found to increase from 12.53 GPa to 17.83 GPa (Fig. 12(b)). Further, the 4 × 5 × 50 mm long bars of SM4 composites were welded to themselves using a filler of SM1. The appearance of weld is similar to that of metal welds (Fig. 13(a)), and is very clean and free from cracks and oxidation of either parent material or filler material. The formation of cracks and pores on either side of the joint interface in heat affected zone was observed after welding without filler. The SM4 welds were subjected to shear testing in steel fixture in shear testing set-up. The joint failed in weld zone. The morphology of fracture surface revealed cleavage/brittle mode fracture. The calculated shear strength was 100 MPa which is lower than the three point bend flexural strength 213 MPa of SM4 composite. The present results revealed the possibility of joining of (ZrB2-SiC-B4C-YAG) composites by GTAW without preheating and post-controlled cooling. Thus in this study oxidation resistant, thermal shock resistant, easily formable/ weld able (ZrB2-SiC-B4C-YAG) composites could be pressureless sintered at relatively low temperatures between 1550 and 1680 °C.  4. Conclusions  Pressureless sintering of (ZrB2-SiC-B4C) composites with (Y2O3 + Al2O3) additions has been studied by changing the vol.% of SiC, B4C, and (Y2O3 + Al2O3) from (26, 24, and 16) to (5, 4, and 5) respectively. A decrease in hardness and increase in flexural strength have been observed with a decrease in B4C, and (Y2O3 + Al2O3) respectively. Composite with high vol.% of SiC, B4C, and (Y2O3 + Al2O3) has been found to retain room temperature flexural strength up to a temperature of 1500 °C. The composite with 20 vol.% SiC, 8 vol.% B4C, and 7 vol.% (Y2O3 + Al2O3) possessing a density of 4.64 g cm− 3, flexural strength of 213 MPa, and Vickers hardness of 17.3 GPa has been sintered. Conical shapes could be made by manual shaping and sintering. A complex yttria-alumina-silicate layer was found to protect  Fig. 11. EDS of (a) oxide layer of HCP ZrB2 in Fig. 10(a), and EDS of (b) bright and gray phases in oxide layer of SM1 in Fig. 10(c).  modulus, strength, fracture toughness, thermal expansion, thermal conductivity, and rate of heat transfer [52]. High thermal expansion coefficient (5.9 × 10 − 6 K− 1), high Young's modulus (489 GPa), and low fracture toughness (3.5 MPa m− 2), result in a thermal shock failure of ZrB2 through crack initiation during rapid heating or cooling, or in the presence of large thermal gradients [53]. Thermal shock failure through crack initiation can be decreased by reducing Young's modulus (E) and co-efficient of thermal expansion (CTE) by pre-heating. Young's modulus of ZrB2 has been reported to decrease by 50% from 524 GPa at room temperature to 263 GPa at 1600 °C. This change in E is equal to approximately doubling the crack initiation thermal shock resistance [54]. Derek S King et al. [31] reported the plasma arc welding of TiB2-20 vol.% TiC composites and ZrB2-20 vol.% ZrC composites [32] by pre-heating of weld coupons. Even after pre-heating to a temperature of 1450 °C and controlled cooling after welding, the formation of porosity is observed at weld interface. Molten liquid of ZrB2 composite will form upon striking an arc between tungsten electrode and the solid surfaces of the composite. As the molten pool is cooled the solid joint surfaces are bonded together. Due to shrinkage accompanied by the solidification of melt pool formation of some porosity at solid-liquid boundary between parent material and fusion zone is possible. During welding, the oxidation of the parent  a   b   50 µm  100 µm   Fig. 12. Optical photo-micrographs of cross section of SM4 weld (a) clean interface marked by arrow, and (b) Vickers micro-hardness across the interface.  \\x0c\", '64  R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  a   b   50 µm  Fig. 13. (a) The appearance of weld of SM4, (b) weld interface, right side SM1 filler material.  the composite from oxidation at high temperature of 1700 °C. The composites exhibited good dimensional stability and thermal shock resistance at 2200 °C in oxy-acetylene flame and at 2700 °C in plasma flame. The formation of yttria stabilized zirconia embedded in the matrix of YAG has been identified on the flame exposed surfaces. The composite could be joined to itself by gas tungsten arc welding with a filler material containing (ZrB2-SiC-B4C-YAG). The shear strength of the weld was about 50% of the flexural strength of the parent composite.  Acknowledgments  Authors thankfully acknowledge the financial support (DMR-295) from the Defence Research and Development Organization, Ministry of Defence, Govt. of India, New Delhi in order to carry out the present study. 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[41]  [35]  [36]  [37]  [46]  [47]  [33]  [39]  [43]  [45]  \\x0c', 'R.V. Krishnarao et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 55-65  65  [49]  P.A. Williams, R. Sakidja, J.H. Perepezko, P. Ritt, Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content, J. Eur. Ceram. Soc. 32 (2012) 3875. [50] X. Zhang, P. Hu, J. Han, S. Meng, Ablation behavior of ZrB2-SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions, Comput. Sci. Technol. 68 (2008) 1718-1726. F. Monteverde, A. Bellosi, Development and characterization of metal-diboridebased composites toughened with ultra-fine SiC particulates, Solid State Sci. 7 (2005) 622-630.  [51]  [52] D. Munz, T. Fett, Thermal shock behaviour, in: D. Muntz, T. Fett (Eds.), Ceramics—Mechanical Properties, Failure Behaviour, Materials Selection, Springer Verlag, Berlin 1999, pp. 203-226. J.W. Zimmermann, G.E. Hilmas, W.G. 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},{
  "_id": 214,
  "PDF": "Pressureless sintering of HfB2–SiC ceramics doped with WC.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 32 (2012) 3627-3635  Pressureless sintering of HfB2-SiC ceramics doped with WC  De-Wei Ni, Ji-Xuan Liu, Guo-Jun Zhang  ∗  State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai  Received 16 March 2012; received in revised form 28 April 2012; accepted 1 May 2012  Available online 23 May 2012  200050, China  Abstract  Using WC as sintering aid, nearly full dense (99%) HfB2 -20 vol% SiC ceramics were sintered at 2200 C for 2 h without external pressure. The densification mechanism, microstructure evolution, mechanical properties and oxidation resistance were  investigated. The results  indicated  that complex chemical reactions of WC in HfB2 -SiC system strongly related to the densification, microstructure and properties. The Young’s modulus, toughness and 3-pt bending strength of HfB2 -20 vol% SiC with 10 wt% WC were 511 GPa, 4.85 Mpa m1/2 and 563 MPa, respectively, fracture  which were comparable to some hot pressed HfB2 -SiC ceramics in literature. The oxidation of HfB2 -20 vol% SiC with 10 wt% WC at 1500 C in C for 10 h, its weight gain and SiC-depleted layer thickness were 3.7 mg/cm2 and 43  \\u242em, air exhibited parabolic kinetics. After oxidation at 1500 respectively, and its residual flexural strength was comparable to or even a little higher than the value before oxidation. © 2012 Elsevier Ltd. All rights reserved.           Keywords: Borides; Sintering; Mechanical properties; Oxidation resistance  1.   Introduction     Hafnium diboride is a Group IV metal diboride which belongs to  the  family  of  ultra-high-temperature  ceramics  (UHTCs) known  for  their  combination of  thermo-physical properties, such as melting point (3380 C), Young’s modulus (480 GPa), −1 K −1 ) and hardness (28 GPa), thermal conductivity (104 W m resistance  to  chemical  attack.1-4 As  a  result, HfB2 -based composite has been considered as potential candidate for ultrahigh  temperature  applications  such  as  propulsion  systems, rocket nozzles, sharp  leading edges and nosecones where  the C.1 With  operating  temperature  can  exceed 2000 the  addition of  silicon  carbide, oxidation  resistance  and mechanical properties of HfB2 ceramics  improve significantly. Therefore, HfB2 -SiC system has been drawing great attention  in  the past years.5,6 Due  to  its  strong covalent bonding and  low  self-diffusion coefficient,  the densification of HfB2 -based ceramics was  typically accomplished by pressure-assisted sintering, such as hot pressing, spark plasma sintering, and reactive hot pressing.7-10 Compared with pressure-assisted sintering, pressureless sinter    ∗  Corresponding author.  E-mail address: gjzhang@mail.sic.ac.cn (G.-J. Zhang).  0955-2219/$ - see front matter © 2012 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2012.05.001     ing  allows  the  fabrication of  components  to near-net  shape. However, because of poor sinterability due  to oxygen contamination of non-oxide  raw materials, pressureless  sintering of borides has been proved very difficult. Using 5-20 vol% MoSi2 as sintering aid, Sciti et al. realized the densification of HfB2 at 1900-1950 C via pressureless sintering.11,12 They argued that MoSi2 softens at high temperature, which can fill voids among matrix grains, favor the matter transport and then accelerate the overall densification of HfB2 ceramics. However,  the  softening of MoSi2 at high temperature might limit the usefulness of these ceramics at high temperature. View the progress on pressureless sintering of ZrB2 -based ceramics,  refractory ceramic aids,  such as C, B4C, WC, can activate  the  raw powders by removing  the oxides on  the surface of ZrB2 , and  thus  improve the densification behavior.13-17 As refractory ceramic aids, they could retain  the properties of ZrB2 at elevated  temperature. It has been  reported  that  the bending strength of ZrB2 -20 vol% SiC composite with addition of 5 vol% WC was as high as 675   33 MPa at 1600 C, which was even a  little higher  than temperature.18 Recently,  the value at room  the densification of HfB2 was also realized via pressureless sintering using refractory ceramic aids.19,20 Using pre-sintering heat treatment, nearly full dense HfB2 (98.3%) was pressureless sintered at 2200 C with 2 wt% B4C as sintering aid.19 The hardness, Young’s modulus and bending strength of HfB2-2 wt% B4C were 19.5 GPa,  ±              \\x0c', '3628   D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635           529 GPa and 469 MPa, respectively. Further  investigation  indicated the main pressureless sintering mechanism of HfB2-2 wt% B4C ceramic switched from solid state sintering (below 2300 C) to liquid phase sintering (above 2300 C).20 WC is another refractory ceramic aid, which has also realized the pressureless sintering of ZrB2 -based ceramics.18 Furthermore, it has been reported that the addition of ‘W’ can improve the oxidation resistance of ZrB2-SiC ceramics.1 Carney et al.21 also  reported  that  the addition of WB/WC  improved  the oxidation  resistance  of HfB2-SiC  at  temperatures  higher  than 1800 C. Based on  the authors’ knowledge, no data has been published related to pressureless sintering of HfB2-SiC ceramics. Recently, phase pure hafnium diboride powder with  low oxygen content has been  synthesized by borothermal  reduction  and  showed high  sinterability  in hot pressing.22 In  the present study,  the main objectives were  to  identify  the effect of WC on  the densification behavior, microstructure evolution and mechanical properties of HfB2-SiC composite processed by  pressureless  sintering.  The  oxidation  resistance  of  the as-sintered HfB2-SiC-WC composites were evaluated and discussed as well.  2. Experimental procedure  2.1. Powder processing and sintering  x         1.37  \\u242em, purity 98.5%, HfB2 powder (average particle size  O: 0.79 wt%, C: 0.10 wt%, synthesized at 1600 C using HfO2 and amorphous B as reactants),22 ␣-SiC (D50 = 0.45  \\u242em; purity 98.5%, Changle Xinyuan Carborundum Micropowder Co. Ltd., China)  and WC  (average particle  size < 1  \\u242em, purity >99%, Zhuzhou Hard Alloy Co. Ltd., China) were used as  the starting powders.  In  this  study, 0, 1, 3, 5, and 10 wt% WC were added  to  80 vol% HfB2 -20 vol%  SiC  (shortly  for HSWx, where  is  the  content of WC  added),  and were planetary ball-milled  in a polymer  jar using SiC media at a  speed of 560 rpm  for 8 h  in  acetone. After milling,  the  slurries were dried  using  a  rotary  evaporator  at  65 C,  then  the  resulting  powders were  crushed  and  sieved  through  200-mesh screen. Cylindrical pellets with dimensions of  Φ45 mm   15 mm were  shaped  by  uniaxial  pressing  at  30 MPa,  followed  by cold  isostatic  pressing  (CIP)  at  250 MPa  for  5 min. Sintering was carried out  in a high  temperature graphite  resistant furnace  (MRF3338, Materials Research Furnaces  Inc., Suncook, USA)  at  temperatures  ranging  from 2000  to 2200 C for 2 h. From  room  temperature  to 1650 C,  a heating  rate −1 was  of  50 C min used,  and  it  decreased  to  30 C min above  1650 C. Similar  to  the  pre-sintering  heat  treatment in  reference,19 the  total heating  schedule  in  this work  also included other  two  isothermal holds at 1650 C and 2000 C for  0.5 h,  respectively,  in  addition  to  the final  dwelling  of 2 h  at  the desired  sintering  temperatures between 2000  and 2200 C. Typically,  the furnace was heated  in vacuum at  temperatures below 2000 C; and the atmosphere was then switched to flowing argon  (purity 99.99%) when  temperature  reached 2000 C.  −1  ×                                a   2.2. Material’s characterization and oxidation tests  2.2.1. Microstructure  The bulk density of the sintered specimens was measured by Archimedes method using water as immersing medium. The theoretical densities of  the samples were calculated with  the rule of mixture based on  the  initial composition. Phase composition was determined by X-ray diffraction (XRD, D/max 2550 V, Rigaku, Tokyo, Japan) using Cu K␣ radiation. Microstructures were characterized using scanning electron microscopy (SEM, JEOL JXA-8100F, Tokyo, Japan) along with energy-dispersive spectroscopy (EDS; Oxford Instruments, UK) for chemical analysis. The grain size of  the as-sintered ceramics was quantified with an average of 100 grains using an  image analysis  software package  (Image-Pro). The oxygen content of as-milled powder mixtures and part of sintered pellets was measured by Nitrogen/Oxygen Exterminator (TC600, Leco Corporation, St. Joseph, MI).  2.2.2. Mechanical properties  The  Vickers’  hardness  and  fracture  toughness  were determined  by  the  indentation  method  (Wilson-Wolpert Tukon2100B, Instron, Norwood, MA) using a load of 5 kg and a dwell time of 10 s; the reported value was an average of 3 measurements. The fracture  toughness (KIC ) was calculated using the following equation23 :  (cid:2)  (cid:3)  KIC =  P  π  C1 +  C2  4  (cid:4)(cid:5)−3/2  −1  (tan   β)  (1)  β  ×  where P is the indentation load (N), C1 and C2 is the measured diagonal crack  length  (m), and  is an angle constant  (68 ). Flexural strength was measured by a 3-pt bending test (test bars 3 mm   4 mm   36 mm) with a span of 30 mm and the crosshead speed was 0.5 mm/min: the reported strength value is an average of 5 measurements. The Young’s modulus (E) was also obtained from the bending test using the following equation:  ×     =  E  L3 (P2 − P1 ) 4BH 3 (Yt2 −  Yt1 )  (2)  where P1 and P2 are the initial and final load (N) of linear range, respectively, L  is  the span (mm), B and H are  the width (mm) and thickness (mm) of samples, respectively, Yt1 and Yt2 are the deflection (mm) when load are P1 and P2 , respectively.  2.2.3. Oxidation resistance  ×  ×  Bars with dimensions of 3 mm   4 mm   36 mm were cut from  the sintered billets and ground  to a 0.5  \\u242em surface finish for oxidation testing. Oxidation studies were conducted in a box furnace by exposing  the specimens  in stagnant air at 1500 C −1 and natural for 0.5 h, 3 h and 10 h. A heating rate of 10 C min cooling was used. The sample was placed on a zirconia plate with minimal contact area  to  limit  interaction at high  temperature. Mass of the specimens was measured using a balance with 0.1 mg precision before and after oxidation. The microstructural modifications on  the oxidized specimens were evaluated by SEM-EDS observing  the external  surfaces as well as  the                    \\x0c', 'D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635   3629        on the removal of oxygen contamination, the oxygen content of HSW10 pellets at two specific points was also evaluated. After the first holding at 1650 C, the oxygen content of HSW10 was 0.23 wt% and it decreased to  0.07 wt% after the second holding at 2000 C. However, the oxygen content of HfB2-20 vol% SiC composite without WC addition was as high as 0.39 wt% even after the final holding at 2200 C. Therefore, it could conclude  that  the addition of WC actually accelerate  the oxygen contamination removement (its reaction mechanism will be discussed in the following part). The use of two isothermal holdings at 1650 and 2000 C prior to the final sintering temperature also proved  to be very effective for  the oxygen  impurities removal. Besides the effect of removing oxygen impurities, the complex reaction of WC in HfB2-SiC system also strongly related to the densification, which is discussed in the following part.        3.2. Microstructure and phase composition  Fig. 1. Relative densities of HfB2 -20 vol% SiC composites with different content of WC pressureless sintered at 2000-2200 C.     corresponding cross sections. The 3-pt residual flexural strength of  the oxidized  test bars was measured at  room  temperature; the reported value was an average of 3 measurements.  3. Results and discussion  3.1. Densification behavior        The relative densities of HfB2-20 vol% SiC composites with different content of WC pressureless sintered at 2000-2200 C were shown  in Fig. 1. As  the calculation of  theoretical densities were on  the basis of  the  initial composition,  the  relative densities showed here were not precise. However,  it could still reflect  the variational  trend of densities by  combining with their microstructures.  It was  found  that  the densification of HfB2 -20 vol% SiC  just by  increasing  temperature was difficult. The  70% when sintered at 2200 relative density of HfB2-20 vol% SiC without WC was only  C  for 2 h. The addition of WC promoted  the densification of HfB2-20 vol% SiC ceramics effectively. When WC addition was lower than 5 wt%, the densification of HfB2 -20 vol% SiC composites  improved significantly as  the  increase of WC content. Then,  the relative densities had no obvious change when WC content was higher than 5 wt%  in  the whole  temperature range 2000-2200 C. As well,  the densification was accelerated markedly by  increasing the sintering temperature. HfB2-20 vol% SiC-10 wt% WC with relative density higher than 99% could be obtained at 2200 C, which was about 200 C higher than that required for pressureless  sintering of ZrB2-20 vol% SiC-10 vol% WC composite, resulting from the more pronounced refractoriness of HfB2 .16 The chemical analysis of  the ball-milled powder mixture 1.09 wt%. Due to high oxyrevealed the oxygen content were  gen  impurities,  the densification of HfB2 -SiC was completely low relative density (70%) after sininhibited, resulting  in a  tered at 2200 C for 2 h. Generally speaking, WC can react and remove  the oxygen contamination, and  therefore  improve  the densification of HfB2-20 vol% SiC. To study the effect of WC                       Fig. 2 shows  the  typical microstructural  features of HSW5 and HSW10 sintered at 2200 C. Fig. 2(a) and  (b)  is  the secondary  electron  images  (SEI) of  the polished  surfaces  after thermal etching at 1900 C for 10 min. No evident pores could be  found, which  is  in good accordance with  the density measurements. Two main phases can be individuated from the SEM images: the grey phase is HfB2 and the black phase is SiC. At the same sintering conditions, no significant changes of microstructure can be observed as the content of WC increase. HfB2 grains show an equiaxed shape with grain size of 7.4  \\u242em and 6.9  \\u242em for HSW5 and HSW10,  respectively. However, SiC partially developed  to platelet morphology at  temperature of 2200 C, \\u242em and  whose  thickness was about 1-2  lengthwise  size was about 3-8  \\u242em. Resulting from such coarse microstructure,  the main fracture mode of both HSW5 and HSW10 was transgranular (Fig. 2(c) and (d)). But it is certain to still exist proportional intergranular fracture in HSW10, which allowed predicting that HSW10 should have higher fracture toughness. The phase composition of HSW ceramics was determined by XRD and SEM-EDS together. The XRD patterns of HSW10 after dwelling at 1650, 2000 and 2200 C are shown  in Fig. 3. In addition to the main phase HfB2 and SiC, WB, HfC and WC were also detected from the XRD patterns. As sintering temperature increases, the XRD diffraction intensity of WC decreases gradually and disappears at  temperatures higher  than 2000 C. Note  that  the main diffraction peak of SiC  locates at 35.7 is very close to that of WC. And it should point out that the diffraction peak appears at  that position for  the composite sintered at higher  than 2000 C corresponds with  the main  reflection of SiC.  It  is also clear  that  the diffraction peaks of HfC show a shift to higher 2θ  value compared to the corresponding peaks of pure HfC (JCPDS 73-0475). Furthermore, the peak shift seemed more obvious as  temperature  increases. Therefore,  it was considered  that  the “HfC” showed  in  the above XRD of HSW10 was not pure HfC, which should be a solid solution HfC with something else. In this study, W (1.30 ˚A) has a smaller covalent (1.42 ˚A)  radius compared  to Hf  (The covalent  radius of both W and Hf were got from  the periodic  table of elements). If W entered  into  the HfC  lattice,  its smaller covalent radius would             \\x0c', '3630   D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635  Fig. 2. SEM images of polished and fracture surfaces of pressureless sintered HSW5 (a, c) and HSW10 (b, d) composites at 2200     C.  reduce the average unit-cell size of HfC, and shift the peaks of HfC toward high diffraction angle. Hence, we assumed that the “HfC” phase here was HfC containing W. The formation of HfCW solid solution is consistent with the C-Hf-W phase diagram, time ago.24 In addition,  which has been reported  long  there  is an unknown peak at 2θ  = 28.2 for HSW10 sintered at 1650 C, which disappeared when  temperature higher  than 2000 C.  It is preliminarily presumed  that  it was a  transitional phase and detailed study is ongoing.           Fig. 3. XRD patterns of HSW10 composites after dwelling at 1650, 2000 and     2200  C.  As shown  in Fig. 4, four phases detected from  the BEI pictures and corresponding EDS analysis in the as-sintered HSW10 composite, support  the above XRD results (phase composition and solid solution). Except  the expected HfB2 and SiC (1 and 2 in Fig. 4(a), respectively), the brightest phase (3 in Fig. 4(a)) that was  identified as WB by XRD actually was a solid solution of (W, Hf)B. The amount of Hf in (Hf, W)B was less than 0.1 mol%, resulting that peak shift of WB was negligible in the XRD patterns. The relatively darker phase (4 in Fig. 4(a)) was the solid solution of (W, Hf)C as above mentioned in XRD results. Nearly 11 mol% W entered into the HfC lattice, which produced a noticeable peak shift of HfC (Fig. 3). Although the EDS element analysis is not quantitative, we can still conclude that (W, Hf)B and (W, Hf)C solid solution was present in HSW10 along with HfB2 and SiC by combining with the XRD results.  3.3. Chemical reactions and densification mechanisms  Oxygen contamination  (including B2O3 and HfO2 )  is one of  the major factors  inhibiting  the densification of HfB2 -based ceramics. Hence, oxide removal could activate the powders sinterability and accelerate densification. B2O3 easily vaporizes at temperature higher than 1100 C. However, HfO2 could only be removed with the aid of sintering additives. It has been confirmed that WC reacts with and removes the ZrO2 from the surface of ZrB2 particles based on the following reaction16-18 :     3WC   +   ZrO2 →   ZrC   +   3W   +   2CO(g)   (3)  \\x0c', 'D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635   3631  Fig. 4. Typical backscattered electrons image (BEI) and corresponding EDS analysis of HSW10 composite pressureless sintered at 2100-2200     C.  Similarly, we assumed  that WC could also  remove HfO2 according to analogous reaction:  react with and  3WC   +   HfO2 →   HfC   +   3W   +   2CO(g)   (4)  an exchange  follows: HfB2 +   2WC   reaction between HfB2 and WC might occur as  →   HfC   +   2WB   +   C   (5)           Thermodynamic calculations (Fig. 5(a)) indicated that reac2050 tion (4) became favorable above  C in standard state. In in vacuum (10 Pa) below this work, sintering was conducted  2000 C, which decreases  the  favorable  temperature of  reac1300 tion  (4)  to about  C. To better understand  the  reaction between HfO2 and WC, a commercial software package (HSC 5.0, Outokumpu Research Oy, Pori, Finland) was used to predict possible reactions and equilibrium products. The molar ratios of WC-HfO2 and pressure were set as 3:1 and 10 Pa, respectively. The calculated results are shown in Fig. 5(b). According to equation (4), HfC, CO and W were predicted as the final products: the reaction between HfO2 and WC completed at around 1600 C. In this study, the oxygen content in HSW system was about 1.09 wt% after ball milling. Based on the stoichiometry of reaction (4) and densification result (Fig. 1),  5 wt% WC was enough  it could preliminarily conclude  that  to  remove  the HfO2 impurity. For HSW10, it was clear that there was some residual WC after dwelling at 1650 C (Fig. 3). Therefore, there must be other routes to exhaust the remaining WC in HSW10. Based on the reaction products of WB and HfC (actually their solid solution) as identified, here we accordingly supposed that similar to the reaction between TiC and ZrB2 to form ZrC and TiB2 ,25,26           However, the calculation of Gibbs free energy change (\\x01G) 2550 for reaction (5) indicated that it could not occur below  C (Fig. 5(a)). It is noticed that carbon appears as a product in reaction (5) and carbothermal reduction of HfO2 (reaction (6)) has a very large, negative  \\x01G at lower temperature and atmosphere pressure.  If  reactions  (5) and  (6) were merged,  the  favorable reaction temperature between HfB2 and WC (reaction (7), HfO2 involved) would be decreased remarkably. HfO2 +  3C  3HfB2 +   HfC   2CO(g)   HfO2 →   2CO(g)    4HfC    6WB    6WC   →  +  +  (6)  (7)  +  +     Thermodynamic calculations confirmed  the above speculations. The  incorporation of HfO2 made  the  reaction between 1900 HfB2 and WC (reaction (7)) became favorable above  C in standard state. Furthermore, a mild vacuum about 10 Pa was maintained during  the first  sintering  stage  (below 2000 C), which could decrease the favorable temperature of reaction (7) 1300 to  C at 10 Pa (Fig. 5(a)). However, the vacuum will not affect  the reaction  temperature of reaction (5) because  there  is no gas involved. Based on the above analysis, it could be concluded that there were mainly  two different  reactions  in HSW  system. Firstly,        \\x0c', '3632   D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635  Fig. 5.   (a) Calculation of Gibbs   free energy change   (\\x01G)   for   reactions   (4),   calculated by reaction of 3 mol WC and 1 mol HfO2 as a function of temperatures at   (5) and  (7) as a  10 Pa.  function of   temperature;   (b)   thermodynamic equilibrium products  WC could react with HfO2 impurity from the surface of starting HfB2 particles as indicated by reaction (4). The removal of HfO2 promoted  the densification of HSW composites. At  the same time,  ternary reaction among WC, HfB2 and HfO2 could also take place according  to  reaction  (7)  if HfO2 appeared around HfB2 and WC grains. HfO2 , the well known oxide impurity in HfB2 , promoted  the  reaction between WC and HfB2 . The asproduced W  from  reaction  (4) could be  incorporated  into  the HfB2 or HfC  lattice and form  the solid solutions of (W, Hf)B and (Hf, W)C, respectively, which has been confirmed by XRD patterns and EDS results.  3.4. Mechanical properties     As discussed in part 3.1, HSW composites sintered at 2200 C reached relative densities higher  than 99% when WC addition was higher  than 3 wt%. Hence,  the mechanical properties of HSW5 and HSW10 sintered at 2200 C are  listed  in Table 1, where the densities and grain size are also included. The as-sintered HSW5 and HSW10 had Vickers hardness (Hv5) of 13.9   0.2 GPa and 14.6   0.9 GPa, respectively, which was much  lower  than  that of hot pressed HfB2-20 vol% SiC (20 GPa) partly due to its large grain size.22 On the other side, it has been reported that thermal expansion coefficient difference  ±  ±     Table 1  Typical physical and mechanical properties of pressureless sintered HSW5 and  HSW10 composites at 2200     Properties   −3 )  Bulk density (g cm Apparent porosity (%)   −3 )  Theoretical density (g cm  Relative density (%)  HfB2 grain size (\\u242em)  Hv5 (GPa)   E (GPa)   KIC (MPa m1/2 )   σ   (MPa)  C.  HSW5   9.79   0.11   9.91   HSW10  10.07  0.09  10.16  98.9   7.4   13.9   516   3.36   547   ± ± ± ± ±   0.3    0.2    24    0.2    58    0.6   0.9   19  99.1  6.9   14.6   511   4.85   563   ± ± ± ± ±   0.65   46  of various phases could resulted in large residual stresses, further formed microcracks in the interfaces, which decreased the hardness remarkably.27,28 In  this work, mainly four phases existed in HSW composites, the effect of residual stresses resulted from thermal expansion coefficient on  the  low hardness should be non-ignorable. In terms of absolute value, the Young’s modulus (E) of HSW5 and HSW10 ceramics were 516 GPa and 511 GPa, respectively, which were comparable to the reported value (529 GPa) of pressureless sintered HfB2 using B4C as sintering aid.19 E  is very sensitive to the porosity, and a little decrease of relative density can result in remarkable loss of elastic modulus. Therefore, the high Young’s modulus  in  this work associated with  their high relative densities. The bending strength of HSW5 and HSW10 composites were 547 MPa and 563 MPa, respectively, which are much higher than the reported value (469 MPa) of pressureless sintered HfB2 using B4C as sintering aid.19 The fracture  toughness for HSW5 and HSW10 composites were 3.36 MPa m1/2 and 4.85 MPa m1/2 , respectively. It preliminarily considered that higher WC content increased the amount of minor phases ((W, Hf)B and (Hf, W)C solid solutions). Generally, these minor phases could partially accumulate at the grain boundary and weaken it to some extent. Accordingly, there was still proportional intergranular fracture in HSW10 (Fig. 2(d)). As shown in Fig. 6, the Vickers-indentation induced crack propagation path on polished surface of HSW10 composite was observed by SEM, where we could get some useful information to understand its toughening mechanism. Crack deflection and bridging increased energy dissipation during crack propagation, resulting in higher fracture toughness for HSW10 ceramic. On the whole, the mechanical properties in this work were comparable to that of pressureless sintered HfB2 ceramics and some hot pressed HfB2 -SiC ceramics in literature.7,29  3.5. Oxidation resistance  Oxidation resistance is a very important property for UHTCs. Here, we studied  the oxidation  resistance of HSW10 sintered  \\x0c', 'D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635   3633  To better understand the mechanism, we also investigated the morphology and element composition of  the oxidized surface. As shown  in Fig. 8,  it could be seen  that  the oxidized surface was a dense glass  layer, and some white  (grey) particles dispersed  in  the glass. EDS analysis  indicated  that  these particles were HfO2 with some W, Fe, Cr  impurities. For  the oxidation of ZrB2-SiC system,  it has been confirmed  that ZrO2 first dissolved  into  the SiO2-B2O3 liquid  to form a ZrO2 -SiO2-B2O3 (BSZ) liquid at the inner side of the glassy layer. This BSZ liquid would then flow toward the top of the glassy layer as oxidation proceeds. When  the B2O3 was  lost by evaporation at  the outer surface, ZrO2 precipitated from the BSZ liquid.31 In this study, the “HfO2 ” particles were embedded  in  the glass scale. At  the same time, phase separation of the liquid/glass was evident. For a similar field-assisted sintered HSW system, phase separation was not observed until 1800 C in Carney et al.’s results, which might due to its low WC content (3 mol%).21 As  is known, W  is a  transition metal with high cation field strength. The addition of W can alter  the composition of  the outer oxidation glassy layer, resulting in phase separation of the liquid/glass. Glass compositions  that exhibit phase separation posses an increased viscosity in the two-phase region compared with a single-phase  liquid, further decelerates  the diffusion of oxygen and improves the oxidation resistance. As oxidation temperature  increases (>1800 C),  the outer borosilicate glass will vaporize, resulting the HfB2-SiC ceramic exposes to air directly. Nevertheless, the addition of W can overcome this issue to some extent. The HfO2-WO3 phase diagram shows a solubility of WO3 of about 5 mol% at 1600-2000 C, while above 5 mol%           Fig. 6. SEM   image of Vickers-indentation   induced crack propagation path on  polished surface of HSW10 composite pressureless sintered at 2200  C.              at 2200 C. HSW10 composite oxidized at 1500 C produced oxide scales similar to the reported HfB2-20 vol% SiC.6 Fig. 7 showed  the cross-sectional  image and compositional maps of polished HSW10 specimen after oxidation at 1500 C for 3 h. The oxide scales are composed of an outer fully dense protective layer of SiO2 -rich glass and an underneath porous SiC-depleted layer (mainly consisted of HfO2 and HfB2 , some residual SiC was  included as well). The  thickness of  the above mentioned 18  oxidation scales is  \\u242em and 26  \\u242em, respectively. According to previous  result,  the outer glass  layer also contained  some B2O3 and HfO2 particles,  the surface glassy  layer should be a borosilicate glass, composed of HfO2-SiO2-B2O3 .30  Fig. 7. SEM image and element mappings of polished cross-section of HSW10 composite after oxidation at 1500     C for 3 h.  \\x0c', '3634   D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635  Fig. 8. Surface morphology and EDS pattern of HSW10 composite after oxidation at 1500     C for 3 h.  oxidation time at 1500  C on the room-temperature flexural strength of HSW10.  thickness of SiC-depleted   layer as a function of oxidation   time for HSW10 composite at 1500     C; (b) effect of  Fig. 9.   (a) Weight gain per unit surface area and         the HfO2 solid solution will be accompanied by a liquid.32 The resulting liquid has a melting point of 1473 C, which can promote  liquid-phase sintering of porous HfO2 and  improve  the oxidation resistance of HSW at high temperature. Usually, the thickness of the whole oxidation scales was used to evaluate  the capability of oxidation resistance. Actually,  the thickness of the SiC-depleted layer was approximately the thickness of the material damage induced by oxidation. Therefore, it was more suitable  to evaluate  the degree of materials damage induced by oxidation using the SiC-depleted layer thickness.31 Combining  the weight gain data with  the SiC-depleted  layer thicknesses  from Fig. 9(a),  the oxidation kinetics of HSW10 was analyzed using a generalized power rate equation31,33 :  xn =  kt  (8)  where x is the change in weight or thickness, n is the exponent, k is the rate constant and t is the oxidation time. The kinetic parameters (n, k) for weight gain and SiC-depleted  layer  thicknesses were  (2, 1.39) and  (2, 179),  respectively. And  the correlation coefficient (R) for both weight gain and SiC-depleted layer thicknesses was 0.997, which indicated that fitting of the kinetic data was very good. The oxidation of HSW10  followed parabolic kinetics no matter using  the weight gain data or SiC-depleted layer thicknesses. After oxidation at 1500 C for 10 h, its weight     and  layer   thickness were 3.7 mg/cm2  gain and SiC-depleted  43  \\u242em, respectively. The influence of oxidation on bending strength of HSW10 is shown in Fig. 9(b). As oxidation proceeded, the residual flexural strength increased sharply at first after 0.5 h oxidation, and then gradually decreased  for  longer  time. The  increase of strength after oxidation was also reported in SiC, Si3N4 and ZrB2-SiC, which was attributed to the formation of a thin, dense oxide layer that could heal the surface flaws resulting from sample processing and machining.31,34,35 However, the flaw healing effect was counterbalanced by the generation of new defects, either within the oxide scale or at  the  interface between  the oxide scale and bulk materials as  the oxide  layer became  thicker.35 Therefore, the residual strength gradually decreased as oxidation time went on. But still, the residual strength of HSW10 after oxidation for 10 h was comparable  to or even a  little higher  than  that of  the un-oxidized specimens due to its good oxidation resistance.  4. Conclusion  HfB2-20 vol% SiC-WC composites were pressureless sintered with WC  as  additives  (0-10 wt%)  at  2000-2200 C. Research  indicates  that HSW composites sintered at 2200 C reached relative densities higher  than 99% when WC addition was higher  than 3 wt%. The addition of WC could remove  the HfO2 impurity and react with HfB2 , resulting the formation of           \\x0c', 'D.-W. Ni et al. / Journal of the European Ceramic Society 32 (2012) 3627-3635   3635  (W, Hf)B and(Hf, W)C  in  the as-sintered composites, which promoted  the pressureless densification of HfB2 -SiC composites. The Young’s modulus (511-516 GPa), fracture  toughness (3.36-4.85 Mpa m1/2 ) and 3-pt bending strength (547-563 MPa) of  the as sintered HSW composites were comparable  to some hot pressed HfB2-SiC ceramics  in  literature. The oxidation of HSW10 followed parabolic kinetics no matter using the weight gain data or SiC-depleted layer thicknesses. After 10 h oxidation at 1500 C, the weight gain and SiC-depleted layer thicknesses were 3.7 mg/cm2 and 43  \\u242em, respectively; its residual strength was comparable  to or even a  little higher  than  that of  the unoxidized specimens. Accordingly, the addition of WC not only realized  the pressureless densification of HfB2-20 vol% SiC composites, but also improved its properties.     Acknowledgements  Financial supports  from  the Chinese Academy of Sciences under  the Program  for Recruiting Outstanding Overseas Chinese (Hundred Talents Program), the National Natural Science Foundation of China  (No. 50632070),  the Science and Technology Commission of Shanghai  (No. 08520707800 and No. 09ZR1435500), and  the CAS Special Grant  for Postgraduate Research, Innovation and Practice are greatly appreciated.  References  1. Fahrenholtz WG, Hilmas GE, Talmy IG, Zaykoski JA. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90:1347-64.  2. Opeka M, Talmy   IG, Wuchina EJ, Zaykoski JA, Causey SJ. Mechanical,  thermal and oxidation properties of refractory hafnium and zirconium compounds. J Eur Ceram Soc 1999;19:2405-14.  3. Kalish D, Clougherty EV, Kreder K. Strength, fracture mode, and  thermal stress resistance of HfB2 and ZrB2 . J Am Ceram Soc 1969;52:30-6. 4. Ni DW, Zhang GJ, Kan YM, Sakka Y. 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Formation of tough interlocking microstructure  in ZrB2 -SiC-based ultrahigh-temperature ceramics by pressureless sintering. J Mater Res 2009;24:2428-34.  16. Zou J, Sun SK, Zhang GJ, Kan YM, Wang PL, Ohji T. Chemical   reac tions, anisotropic grain growth and sintering mechanisms of self-reinforced ZrB2 -SiC doped with WC. J Am Ceram Soc 2011;94:1575-83. 17. Zou J, Zhang GJ, Sun SK, Liu HT, Kan YM, Liu JX, et al. ZrO2 removing reactions of groups IV-VI transition metal carbides in ZrB2 based composites. J Eur Ceram Soc 2011;31:421-7.  18. Zou   J, Zhang GJ, Hu CF, Nishimura T, Sakka Y, Vleugels   J, Van der     C. J Am Ceram Soc  Biest O. Strong ZrB2 -SiC-WC ceramics at 1600 2012;95:874-8.  19. Zou J, Zhang GJ, Kan YM, Ohji T. Pressureless sintering mechanisms and  mechanical properties of hafnium diboride ceramics with pre-sintering heat treatment. Scripta Mater 2010;62:159-62.  20. Zou J, Zhang GJ, Kan YM. Pressureless densification and mechanical prop erties of hafnium diboride doped with B4C:  from solid state sintering  liquid phase sintering. J Eur Ceram Soc 2010;30:2699-705.  to  21. Carney CM, Parthasarathy TA, Cinibulk MK. Oxidation   resistance   of  hafnium diboride ceramics with additions of silicon carbide and  boride or tungsten carbide. J Am Ceram Soc 2011;94:2600-7.  tungsten  22. Ni DW, Zhang GJ, Kan YM, Wang PL. Hot pressed HfB2 and HfB2 -20 vol% SiC ceramics based on HfB2 powder synthesized by borothermal reduction  of HfO2 . Int J Appl Ceram Technol 2010;7:830-6.  23. Evans AG, Charles EA. Fracture toughness determinations by indentation.  J Am Ceram Soc 1976;59:371-2.  24. Rudy E, Hoffman JR. Planseeber Pulvermet 1967;15:174-8.  25. Zhang GJ, Wu WW, Kan YM, Wang PL. Ultra-high temperature ceramics (UHTCs) via reactive sintering. Key Eng Mater 2007;336-338:1159-63.  26. Mishra SK, Pathak LC. 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Kinetics and mechanism of oxidation of silicon nitride bonded silicon carbide ceramic. J Therm Anal Calorim 1994;42:811-22.  34. Kim HW, Kim HE, Song H, Ha J. Effect of oxidation on the room tempera ture flexural strength of reaction-bonded silicon carbides. J Am Ceram Soc 1999;82:1601-4.  35. Kim HW, Kim HE. Oxidation and strength retention of monolithic Si3N4 and nano-composite Si3N4 -SiC with Yb2O3 as a sintering aid. J Am Ceram Soc 1998;81:2130-4.  \\x0c']"
},{
  "_id": 215,
  "PDF": "Pressureless Sintering of Zirconium Diboride Using Boron Carbide and Carbon Additions.pdf",
  "Text": "['Pressureless Sintering of Zirconium Diboride Using Boron Carbide and  Carbon Additions  Sumin Zhu,*,w  William G. Fahrenholtz,* Gregory E. Hilmas,* and Shi C. Zhang*  Department of Materials Science and Engineering, University of Missouri-Rolla, Rolla, Missouri 65409  The synergistic roles of boron carbide and carbon additions in the enhanced densification of zirconium diboride (ZrB2) by pressureless sintering have been studied. ZrB2 was sintered to 499% relative density at 19001C. The combination of 2 wt% boron carbide and 1 wt% carbon promoted densification by  removing surface oxide impurities (ZrO2 and B2O3) and inhibiting grain growth. Four-point bending strength (473743 MPa), (19.670.4 GPa), Vickers’ microhardness fracture toughness (3.570.6 MPa . m1/2), and Young’s modulus were measured. Thermal gravimetry showed that the combina (507  GPa)  tion of additives did not have an adverse effect on the oxidation  behavior.  I.  Introduction  Z IRCONIUM DIBORIDE (ZrB2)  is a covalent compound with a  hexagonal AlB2 structure. Among its properties, ZrB2 has an ultra-high melting point (430001C), high Young’s modulus (B500 GPa), high hardness (23 GPa), (B10\\x005 O \\x01 cm), high thermal conductivity 460 W \\x01 (m \\x01 K)\\x001, low electrical resistivity and resistance to chemical attack.1,2 The combination of these  characteristics makes ZrB2 and ZrB2-based ceramics useful for high-temperature structural applications, such as cutting tools,  thermal protection for hypersonic vehicles, molten metal cruci bles, high-temperature electrodes, and thermowell tubes for steel refining.3,4  Densification of ZrB2, without additives, typically requires hot pressing at 21001-23001C due to its strong covalent bonding and the presence of surface oxide impurities.5 Dense ZrB2 ceramics have been fabricated at temperatures below 20001C with  additives such as metals (e.g., Fe, Ni) or Si3N4).6 Despite improved densification, derived from these additives can cause mechanical degradation  ceramics  (e.g., SiC,  the secondary phases  at high temperatures. Furthermore, hot pressing is  limited to  simple geometries. In comparison, pressureless sintering offers  the advantage of processing to near net shape. Several sintering  additives have recently been reported to promote pressureless  sintering of ZrB2. For example, ZrB2-20 vol% MoSi2 was sintered to near full density at 18501C.7 Enhanced densification was  attributed to liquid-phase  formation. Further, ZrB2 milled using WC media was sintered to near theoretical density at 18501C with the addition of 4 wt% B4C8 or at 19001C with 1.7 wt% carbon.9 Enhanced densification was attributed to the  attrition  submicrometer particle size and the removal of surface oxides  (ZrO2 and B2O3) by reaction with the sintering aids. However, with either B4C or carbon alone, as-received ZrB2 powder (i.e.,  no attrition milling) only reached B95% relative density (RD) after sintering at 20001C. In addition, carbon additions of 1.7  wt% or more resulted in the presence of residual carbon in the sintered ceramics.9 The residual carbon was thought  to be un desirable because it usually reduces  the  strength of nonoxide  ceramics.  In this communication, the densification of ZrB2 powder was improved using a combination of 2 wt% B4C and 1 wt% carbon. Thermodynamic analysis was performed to assess the syn ergistic  roles  of  the  sintering  additives. The microstructure,  mechanical properties, and oxidation behavior of  the sintered  ZrB2 ceramics were characterized.  II.  Experimental Procedure  Commercially available ZrB2 (Grade B, H.C. Starck, Karlsruhe, Germany) and B4C (Grade HS, H.C. Starck) powders were used as raw materials. The main characteristics of these powders are  summarized in Table I. Carbon was introduced in the form of  phenolic resin (Type GP 2074, Georgia Pacific, Atlanta, GA),  which had a carbon yield of 41 wt% as determined by thermo gravimetric analysis (TGA, Model STA409, Netzsch, Selb, Ger many). As-received ZrB2 powder was mixed with 2 wt% B4C powder (4.5 vol%) and 1 wt% carbon (2.5 vol%) in acetone by  ball milling for 24 h using WC media. The slurry was then dried  and granulated by pulverizing with a mortar and pestle and to \\x0080 mesh. The powder was uniaxially pressed at 30 MPa into cylindrical pellets (44-mm diameter and 4-mm  screening  tall)  and  further  compacted  by  cold  isostatic  pressing  at  300 MPa. The pellets were heated at 11C/min in flowing argon to 7001C  to char the phenolic resin and produce carbon. Previous analysis  has shown that addition of carbon by this method resulted in the formation of a thin coating on the particle surfaces.9 After char ring,  sintering was  carried out  in a graphite  furnace  (Model  3060-FP20, Thermal Technology, Santa Rosa, CA). Specimens  were placed in a graphite crucible lined with a boron nitride coated graphite foil and heated from room temperature 16001C in a mild vacuum (B25 Pa). The heating rate was 101C/min with 1 h isothermal holds at 14501 and 16001C. Above 16001C, the furnace was backfilled to 1 atm with flowing argon  to  to reduce grain coarsening caused by vapor transport. The furnace was heated at 301C/min from 16001 to 19001C, followed by a 2 h dwell before cooling to room temperature at 301C/min.  For comparison, the same ZrB2 powder without additives was also sintered using the same procedure.  The bulk density of the sintered specimens was measured by  the Archimedes method, while the relative density was estimated using a true density value (5.9 g/cm3) calculated using the rule of  mixtures. Microstructures were characterized using a field emis sion scanning electron microscope (SEM, Model S4700, Hitachi,  Tokyo, Japan), equipped with energy-dispersive X-ray spectros copy (EDS, Model Phoenix, EDAX, Mahwah, NJ). ZrB2 grain sizes were determined from SEM images using an image analysis  software package (ImageJ, National Institutes of Health, Beth R. Speyer—contributing editor  This work was  supported by the National Science Foundation under grant number  DMR-0346800.  *Member, American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: sz6v5@umr.edu  Manuscript No. 22976. Received March 23, 2007; approved June 11, 2007.  Journal  J. Am. Ceram. Soc., 90 [11] 3660 - 3663 (2007)  DOI: 10.1111/j.1551-2916.2007.01936.x  r 2007 The American Ceramic Society  3660  \\x0c', 'November 2007  Communications of the American Ceramic Society  7ZrO2 þ 5B4C ! 7ZrB2 þ 5COðgÞ þ 3B2O3 ðl Þ  2ZrO2 þ B4C þ 3C ! 2ZrB2 þ 4COðgÞ  ZrO2 þ B2O3 ðl Þ þ 5C ! ZrB2 þ 5COðgÞ  3661  (1)  (2)  (3)  Table I.  Characteristics of the Starting ZrB2 and B4C Powders  Material  Grade  Average  particle  size (mm)  ZrB2 B4C  B  HS  2  0.8  Specific  surface area  Oxygen  (m2/g)  1.0  15.8  content (wt%)  Supplier  0.9  1.3  H.C. Starck  H.C. Starck  esda, MD). The average grain size was estimated by measuring  at  least 100 grains. Thermochemical software (HSC Chemistry  5.11, Outokumpu Research Oy, Pori, Finland) was used for the  thermodynamic  calculations. Young’s modulus  (E) was mea sured on a 35-mm diameter disk-shaped specimen by impulse  excitation of vibration (Model MK4-I Grindosonic, J. W. Lem mens, St. Louis, MO) according to ASTM standard C1259-01.  The strength was measured according to ASTM standard C1161-02c. Type A bars (1.5 mm \\x02 2 mm \\x02 25 mm) were fractured in four-point bending using inner and outer spans of 10  mm and 20 mm, respectively, and a crosshead speed of 0.2 mm/  min. Ten bars were  tested. Vickers’ microhardness was mea sured by indentation on polished sections with a load of 9.8 N  and a dwell time of 15 s (Model Duramin-5, Struers Inc., Westlake, OH). The direct-crack measurement method10 was used to  estimate the fracture toughness (KIc). At least 10 measurements were averaged for the reported microhardness and fracture  toughness values. Oxidation behavior was evaluated by TGA in an alumina crucible at a heating rate of 51C/min in flowing (25 cm3/min) up to 15001C. The test dry air specimen was a rectangular bar with dimensions of 1.5 mm \\x02 2.0 mm by 10 mm and was cut from a sintered pellet.  III.  Results and Discussion  As shown in Table II, ZrB2 sintered at 19001C with a combination of B4C and carbon reached 499% RD. For comparison, additive-free ZrB2 reached only 61% RD at the same temperature. Previous research showed that the same ZrB2 powder sintered to B95% RD with 4 wt% B4C at 20501C8 or B95% RD with 2 wt% carbon at 19001C.9 Therefore, the combined  additives enhanced the densification of ZrB2 more than the individual additives.  The removal of surface oxides has been shown to be critical to  the densification of boron-containing ceramics, such as B4C and TiB2.11,12 It is generally believed that surface oxygen contamination in the form of B2O3 and ZrO2 inhibits the densification of ZrB2 by promoting coarsening through evaporation-condensation at intermediate temperatures (15001-18001C).6,11 The addi tion of reactive additives has been found to enhance densification substantially by removing surface oxygen impurities.8,9  In the  present study, B4C and carbon were used together. Carbon was added in the form of phenolic resin, which enabled it to be dis tributed uniformly around the ZrB2 and B4C through the formation of carbon coatings on the powder surfaces.9 The  additives  reduce  the oxides  through processes  like  those de scribed by the following reactions:  Table II.  Physical and Mechanical Properties of ZrB2 Sintered at 19001C  Property  Relative density  ZrB2 grain size Four-point bend strength  Fracture toughness  Vickers’ microhardness  Young’s modulus  Value  499% 4.172.0 mm 473743 MPa 3.570.6 MPa \\x01 m1/2 19.670.4 GPa  507 GPa  Because of the actively pumped vacuum used for sintering below 16001C, the CO produced as a gaseous product was con stantly removed from the system, and its partial pressure would, (B25 Pa).  therefore, remain below the measured vacuum level  Thermodynamic calculations indicate that reactions (1)-(3) become favorable above 9281, 10001, and 10441C, respectively,  assuming a CO partial pressure of 25 Pa.  For ZrB2, both ZrO2 and B2O3 are present on the surface. Thermodynamically, reaction (1) has the highest driving force  and is thought to be the main reaction responsible for ZrO2 removal when ZrB2 is sintered with B4C alone.8 However, this reaction leads to the formation of liquid B2O3, which will promote grain coarsening (and thereby inhibit densification) if it is removed. Although B2O3 volatilizes above 12001C under vacuum when no reactive additives are used, around 10% liquid  not  B2O3 may be retained in ZrB2 powder compacts up to at least 14001C in the menisci between the grains.13 Analysis of reactions  (2) and (3) indicates that both ZrO2 and liquid B2O3 can be removed by reaction in the presence of carbon and B4C. Thus, the carbon may facilitate B2O3 removal more effectively than when B4C is used alone. Thermodynamically, the combination of B4C and carbon additions removed surface oxides more effectively  compared with either additive alone by reacting with both ZrO2 and B2O3.  Fig. 1. (a) Polished section and (b) fracture surface of ZrB2 sintered at 19001C with a combination of B4C and carbon additions. The gray phase indicates ZrB2, the dark phase indicates B4C, and the residual carbon is not discernible.  \\x0c', '3662  Communications of the American Ceramic Society  Vol. 90, No. 11  Figure 1 shows the microstructure of ZrB2 sintered at 19001C with B4C and carbon additions. Much of the apparent porosity in the polished section is thought to be due to grain pullout  (Fig. 1(a)). SEM observations confirmed that the ceramic reached near full density, with o 1% closed porosity measured by image analysis. EDS analysis indicated that the dark phase  was B4C, which was mainly though a small fraction was located within ZrB2 grains. Residual carbon, however, was not discernible from the SEM or EDS  intergranularly,  distributed  al analysis. Carbon could be consumed either by reaction with ox ides (reactions (2) and/or (3)) or by dissolution into the ZrB2 matrix. The binary ZrB2-C phase diagram14 indicates that the carbon solubility in ZrB2 is o2 at.% (B0.2 wt%), implying that only a small fraction of the carbon could be consumed by dis solution into ZrB2. The fracture surface (Fig. 1(b)), with a higher magnification inset image, was typical of intragranular  fracture, which indicates  strong bonding in the grain bound aries owing to the removal of oxygen contamination. Despite  the fact that WC media were used for ball milling, no WC was  observed in the microstructure, which indicated that WC impu rity content was minimal  in this material.  Table II contains the grain size and mechanical property data for the sintered ZrB2 ceramic. The grain size (4.172.0 mm) of dense ZrB2 was approximately twice as large as the average particle size (2 mm) of the starting powder. Therefore, significant  grain growth was inhibited by the presence of B4C and carbon in the grain boundaries.8,9 The ZrB2 ceramic with a grain size of B4 mm had a four-point bend strength of 473743 MPa, which  is about twice as high as the strengths that have been reported  for ZrB2 ceramics with grain sizes of several tens of micrometers (200-375 MPa).3,15 The finer grain size and oxygen-free grain  boundaries are, presumably, the main reasons for the improved bend strength. The Vickers’ microhardness was 19.670.4 GPa, (B20  which is comparable with that GPa).16 The grain boundaries  of  single-crystal ZrB2 should pin dislocations, which  may result  in higher hardness of polycrystalline ZrB2 than the single crystal. However, hardness of the sintered ceramic can  also be lowered significantly due to the residual stresses caused by the anisotropic thermal expansion behavior of ZrB2.17 The (3.570.6 MPa \\x01 m1/2) was what lower than some values reported for ZrB2 (4-5 MPa \\x01 m1/2), measured fracture toughness some3 but comparable with a value of 3.570.3 MPa \\x01 m1/2 reported recently for the hot-pressed ZrB2 with a grain size of B6 mm.18 The additions of B4C and carbon do not seem to enhance the fracture toughness of ZrB2 ceramics. As shown in Fig. 1(b), fracture was predominantly  intragranular. Thus,  crack the  deflecting mechanisms were  not  operating  and  the  fracture  toughness was  sintered ZrB2 was 507 GPa, which agrees well with the single-crystal value of 500 GPa.19  low. Young’s modulus  (E) of  The mass gain for pure ZrB2 ceramics has been generally reported to be more than 10 mg/cm2 up to 15001C during nonisothermal oxidation in dry air.4,20 As  shown in Fig. 2,  the  sintered ZrB2 in this study showed a normalized mass gain of 8.5 mg/cm2 after heating to 15001C. The lower mass gain may be  due to the high RD. In addition, the presence of B4C could facilitate the formation of a glassy B2O3 layer that would act as a barrier to the diffusion of oxygen.21 Based on Fig. 2, oxidation  of sintered ZrB2 was also found to be comparable with that reported for hot-pressed materials4 despite the addition of two  sintering aids (B4C and carbon).  IV.  Summary  ZrB2 powder was densified by pressureless sintering to 499% RD at 19001C using a combination of B4C and carbon additives. The addition of 2 wt% B4C and 1 wt% carbon not only promoted densification by reacting with the surface oxides but  also inhibited grain growth. Starting from a commercial powder with an average ZrB2 particle size of 2 mm, the sintered ceramic had an average ZrB2 grain size of B4 mm. A fine grain size and a  Fig. 2. Mass gain in air as a function of temperature for sintered ZrB2 by thermogravimetric analysis at a heating rate of 51C/min.  lower oxygen content resulted in a high strength (473743 MPa).  In contrast, the additions did not improve fracture toughness, which was B3.5 MPa \\x01 m1/2. Young’s modulus (507 GPa) and Vickers’ microhardness (19.670.4 GPa) were comparable with  those reported for single-crystal ZrB2. The additions of B4C and carbon did not impair the oxidation resistance of ZrB2. Based on this study, the combination of B4C and carbon promoted the densification of ZrB2 at lower temperatures than either additive alone by reacting with and facilitating the removal of both B2O3 and ZrO2 from the particle surfaces.  References  1R. A. Cutler, ‘‘Engineering Properties of Borides’’; pp. 787-811 in Ceramics and  Glasses, Engineered Materials Handbook, Vol. 4, Edited by S. J. Schneider Jr..  ASM International, Materials Park, OH, 1991. 2M. J. Gasch, D. T. Ellerby, and S. M. Johnson,  ‘‘Ultra High Temperature  Ceramic Composites’’; pp. 197-224 in Handbook of Ceramic Composites, Edited  by N. P. Bansal. Springer, New York, 2005. 3C. Mroz, ‘‘Zirconium Diboride,’’ Am. Ceram. Bull., 74 [6] 164-5 (1995). 4W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. Zaykoski,  ‘‘Refrac tory Diborides of Zirconium and Hafnium,’’ J. Am. Ceram. Soc., 90 [5] 1347-64  (2007). 5H. Pastor,  ‘‘Metallic Borides: Preparation of Solid Bodies. Sintering Methods  and Properties of Solid Bodies’’; pp. 457-93 in Boron and Refractory Borides,  Edited by V. I. Matkovich. Springer, New York, 1977. 6F. Monteverde, S. Guicciardi, and A. Bellosi, ‘‘Advances in Microstructure and  Mechanical Properties of Zirconium Diboride-Based Ceramics,’’ Mater. Sci. Eng.,  A346, 310-9 (2003). 7D. Sciti, S. Guicciardi, A. Bellosi, and G. Pezzotti,  ‘‘Properties of a Pressure less-Sintered ZrB2-MoSi2 Ceramic Composite,’’ J. Am. Ceram. Soc., 89 [7] 2320-2 (2006). 8S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, ‘‘Pressureless Densification  of Zirconium Diboride with Boron Carbide Additions,’’ J. Am. Ceram. Soc., 89 [5]  1544-50 (2006). 9S. Zhu, W. G. Fahrenholtz, G. E. Hilmas, and S. C. Zhang,  ‘‘Pressureless  Sintering of Carbon-Coated Zirconium Diboride Powders,’’ Mater. Sci. Eng. A,  A459, 167-71 (2007). 10G. R. Anstis, P. Chantikul, B. R. Lawn, and D. B. Marshall,  ‘‘A Critical  Evaluation of Indentation Techniques for Measuring Fracture Toughness: I, Di rect Crack Measurements,’’ J. Am. Ceram. Soc., 64 [9] 533-8 (1981). 11H. Lee and R. F. Speyer,  ‘‘Pressureless Sintering of Boron Carbide,’’ J. Am.  Ceram. Soc., 86 [9] 1468-73 (2003). 12S. Baik and P. F. Becher,  ‘‘Effect of Oxygen Contamination on Densification  of TiB2,’’ J. Am. Ceram. Soc., 70 [8] 527-30 (1987). 13I. G. Talmy, J. A. Zaykoski, and M. A. Opeka, ‘‘Properties of Ceramics in the  ZrB2/ZrC/SiC System Prepared by Reactive Processing,’’ Ceram. Eng. Sci. Proc., 19 [3] 105-12 (1998). 14‘‘Figure 8874 (C)’’; p. 205 in Phase Diagrams for Ceramists, Vol. X, Edited by  A. E. McHale. The American Ceramic Society, Westerville, OH, 1994. 15G.-J. Zhang, Z.-Y. Deng, N. Kondo,  J.-F. Yang, and T. Ohji,  ‘‘Reactive  Hot Pressing of ZrB2-SiC Composites,’’ (2000).  J. Am. Ceram. Soc.,  83  [9]  2330-2  \\x0c', 'November 2007  Communications of the American Ceramic Society  3663  16Y. Xuan, C. Chen, and S. Otani,  ‘‘High Temperature Microhardness of ZrB2 Single Crystals,’’ J. Phys. D: Appl. Phys., 35, L98-1 (2002). 17A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas,  ‘‘Pressureless  Sintering of Zirconium Diboride,’’ J. Am. Ceram. Soc., 89 [2] 450-6 (2006). 18A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  Single Crystals of ZrB2 and the Suitability of ZrB2 as a Substrate for GaN Film,’’ J. Appl. Phys., 93 [1] 88-93 (2003). 20F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 21C. A. A. Cairo, M. Florian, M. L. A. Graca, and J. C. Bressiani,  ‘‘Kinetic  in Dry  ‘‘High-Strength Zirconium Diboride-Based Ceramics,’’ J. Am. Ceram. Soc., 87 [6]  Study by TGA of  the Effect of Oxidation Inhibitors for Carbon-Carbon Com 1170-2 (2004). 19N. L. Okamoto, M. Kusakari, K. Tanaka, H. Inui, M. Yamaguchi, and S.  Otani, ‘‘Temperature Dependence of Thermal Expansion and Elastic Constants of  posites,’’ Mater. Sci. Eng., A358, 298-303 (2003).  &  \\x0c']"
},{
  "_id": 216,
  "PDF": "Pressureless sintering of ZrB2–SiC ceramics- the effect of B4C content.pdf",
  "Text": "['Available online at www.sciencedirect.com  Scripta Materialia 60 (2009) 559-562  www.elsevier.com/locate/scriptamat  Pressureless sintering of ZrB2-SiC ceramics: of B4C content  the effect  Hui Zhang,a,b,* Yongjie Yan,a Zhengren Huang,b Xuejian Liua and Dongliang Jianga  aShanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China bGraduate School of  the Chinese Academy of Sciences, Beijing 100039, China  Received 1 November 2008; revised 2 December 2008; accepted 2 December 2008  Available online 13 December 2008  Systematic studies were carried out on the sintering of ZrB2-SiC ceramics without any external pressure being applied. The densities, mechanical properties and microstructures with different contents of B4C were investigated. The results showed that B4C additions were beneficial not only for the sintering process but also for the oxidation resistance. The sintered ZrB2-SiC ceramics with 0.4 wt.% B4C addition had an average strength of 382 ± 24 MPa and a modulus of 405 ± 34 GPa. The hardness and toughness were 13.4 ± 0.4 GPa and 3.7 ± 0.2 MPa m1/2, respectively. The effect of B4C addition on the oxidation resistance of ZrB2-SiC was studied by observing the mass changes and the residual strengths of samples after oxidized treatments.  Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: ZrB2-SiC; Microstructure; Mechanical properties  Ultrahigh-temperature ceramics (UHTCs), such as zirconium diboride (ZrB2), HfB2, ZrC and HfC, are remarkable for their ultrahigh melting points (>3000 °C) and most of them belong to the early transition metal borides and carbides [1-3]. In this family of materials, ZrB2-based UHTCs have an unusual combination of properties, such as high strength, high hardness, high electrical and thermal conductivities, good oxidation and chemical attack resistances, and the lowest theoretic density [4-7]. This makes them attractive for aerospace and other applications that require exposure to extreme thermal and chemical environments, such as those associated with atmospheric re-entry, hypersonic flight and rocket propulsion, high-temperature electrodes and crucibles for molten metal contact [4,8-10]. Owing to their strong covalent bonding as well as their low volume and grain boundary self-diffusion coefficients, ZrB2-based ceramics have typically been densified by hot pressing [5,6,11]. Recently, spark plasma sintering (SPS) has been used to densify diborides [12,13]. Because of its rapid heating rate (\\x18100 °C), grain coarsening is limited during the SPS process, which enables densification without significant grain growth. Densification of borides has also been achieved by reactive hot processing [14,15]. However, pressureless sintering offers a number of advan * Corresponding  author. Address:  Shanghai  Institute  of Ceramics,  Chinese Academy of Sciences, 1295 DingXi Road, Shanghai 200050,  China. E-mail: zhrhuang@mail.sic.ac.cn  tages over hot pressing, SPS and reactive hot pressing, including the ability to fabricate components to nearnet shape using standard powder processing methods. Above all, the diamond machining of the sintered parts can be substantially reduced, which can cut the costs dramatically [16]. On the other hand, it is rather difficult to achieve full density of pure ZrB2-based ceramics by pressureless sintering without any additives due to its poor intrinsic sinterability. Incorporation of additives such as TiB2 [5], Si3N4 [17], SiC [18,19], MoSi2 [27,28], Mo [20], Fe [21], Ni [5,11,22], B4C [5,23] and C [23,24] has been reported to improve the pressureless sintering of ZrB2based ceramics. However, the addition of metallic additives and ceramic additives can also apparently deteriorate the properties of the ceramic materials at hightemperatures [21,22]. Recently, some progress has been made in the pressureless sintering of ZrB2-based UHTCs. Chamberlain et al. [7] sintered ZrB2 to a relative density of \\x1898% without applying any external pressure. However, the introduction of WC (\\x182 vol.%) was thought to form a solid solution with ZrB2 and promote the pressureless sintering process. Zhang et al. [23] densified ZrB2 at 1850 °C by pressureless sintering to nearly full density. Both WC and B4C were used as sintering aids to facilitate the densification process. Zhu et al. [25] used B4C and C as the sintering aids to achieve nearly full-density ZrB2 by the pressureless sintering method. Generally speaking, it was supposed that B4C, WC or C was used to remove the oxides on the surface of the ceramic powders which  1359-6462/$ see front matter Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2008.12.003  \\x0c', '560  H. Zhang et al. / Scripta Materialia 60 (2009) 559-562  inhibited the densification process. In some studies, attrition milling was used to decrease the particle size and increase the specific surface area of the powders, which should also have increased the driving force for densification. Nevertheless, systematic studies about B4C additions on the pressureless sintering of ZrB2-based ceramics have not been reported. Moreover, previous investigations have demonstrated that SiC particulate additions can increase the oxidation resistance of ZrB2 by promoting the formation of silicate-based glass phases that inhibit oxidation at temperatures between 800 and 1700 °C [6,26]. In addition to improving the oxidation resistance, the SiC particulates can also improve the sinterability and inhibit the grain growth of the ZrB2. The purpose of this paper was to study the effect of different amounts of B4C on the pressureless sintering process of ZrB2-SiC ceramics. The relative density, mechanical properties, microstructure and oxidation resistance were investigated as the B4C contents changed. Commercially available powders were used as raw materials in this study. The as-received ZrB2 powders (Dandong Chemical Co., Ltd., Dandong, China) had a purity of >90% (metal basis) and an average particle size of <2.5 lm. SiC powders (a-phase, FCP 15C, SIKA TECH., Lillesand, Norway) had a purity of >99% and an average particle size of <0.5 lm. The B4C powders (Mudanjiang Jingangzuan Boron Carbide Co., Ltd., Mudanjiang, China) had a purity of >97% and an average particle size of <1 lm. Finally, phenolic resin (Shanghai QiNan Adhesive Material Factory, Shanghai, China) was added as a binder for the green parts. Upon heating, it pyrolyzed to 30 wt.% carbon, which also acted as a sintering aid. The commercial ZrB2 powders were attrition-milled (Model 01-HD, Union Process Precision Machinery Co., Ltd., Qingdao, China) in ethanol at 300 rpm for 4 h in a Teflon-coated tank, using cobalt-bonded WC media and a cobalt-bonded WC spindle. After that, a certain amount of attrition-milled ZrB2 powders were mixed with 20 vol.% SiC and 3 wt.% phenolic resin and different B4C additions, and then ball milled for 48 h using SiC milling media in ethanol. After milling, ethanol was removed using rotary evaporation (Rotavapor R-124, Bucchi, Flawil, Germany) at a temperature of 70 °C. The resulting powders were crushed using an agate mortar and pestle and then passed through a 100-mesh sieve to produce uniform granules for dry pressing. The obtained materials contained B4C in amounts of 0, 0.2, 0.4, 0.6, 0.8, 1.0, 2.0 and 4.0 wt.% (based on the weight of the ceramic dry-pressed at \\x1860 MPa using a steel die, followed by cold powders). The composite powders were first uniaxially isostatic pressing at \\x18200 MPa for 120 s. The green compacts were heated to 900 °C and held for 1 h in vacuum to remove the binder at a rate of 5 °C min\\x001. After binder removal, the samples were sintered to 2200 °C and then held for 2 h in flowing argon at a rate of 10 °C min\\x001. The initial vacuum was generally less than 10 Pa and above 1600 °C the furnace was backfilled with argon. After sintering, the furnace was cooled at a rate of 10 °C min\\x001 to room temperature under flowing argon. All the heattreatment processes were carried out in a high-temperature graphite resistance furnace (High-Multi 10000, Fujidempa kogyo, Ltd., Saitama, Japan).  The crystallite size and morphology of the starting powders were observed by field emission scanning electron microscopy (SEM) (JSM-6700F; JEOL, Tokyo, Japan). The microstructure of the sintered body was observed by electron-probe microanalysis (EPMA) (JXA-8100, JEOL) with energy-dispersive spectroscopy (EDS). The bulk density was measured using the Archimedes displacement method with water as the immersing medium and the relative density was calculated by dividing the measured bulk density by a calculated theoretical density. The theoretical density was estimated using ruleof-mixtures calculations that assumed the nominal compositions of the batches as specified. Flexural strength was measured by the three-point bending method with a span of 30 mm and a cross-head speed of 0.5 mm min\\x001, using an Instron 5566 universal testing machine. The size the samples was nominally 3 mm \\x02 4 mm \\x02 36 mm. of Elastic modulus was calculated by the bending method based on the Chinese Standard GB/T 10700-2006. Vickers hardness and fracture toughness were measured by Vickers indentation tests (Model 300, Tukon, Canton, MA) using a load of 10 kg and a dwell time of 10 s. Fracture toughness was calculated by the direct crack measurement method based on the Japanese Standard JIS R 1607-1995. Measurements from at least five indents were averaged to calculate the final value. The particle size and morphology of the ZrB2 powders were observed by scanning electron microscopy in Figure 1. The SEM images in Figure 1(b) revealed that the attrition-milled ZrB2 powders had a flake-like morphology, with an average particle size of about 300- 500 nm. The specific surface area of the milled powder 10m2 g\\x001. Compared with was about the as-received ZrB2 powders in Figure 1(a), attrition milling had broken up the agglomerations to a certain extent and the submicron powders with a high surface area should have improved the sinterability. Elemental analysis of the milled powders revealed increased O, C and W contents (Table 1) compared to the as-received powders. The increased O was thought to be from the formation of the ZrO2 and B2O3 on the exposed surface of the submicron powders. Metallic impurities (W, C) were introduced during milling because of the wear of the WC milling media. The density and mechanical properties of the sintered ZrB2-SiC ceramics are listed in Table 2. The density of the sintered body affected the microstructure and the properties of the materials. From Table 2, as the amount of B4C inclusion increased from 0 to 4 wt.%, the theoretical density of the samples decreased by 4.4% due to the lower density of B4C compared to ZrB2. On the other hand, as with the increased B4C contents from 0 to 0.4 wt.%, the relative density of the sin (a) as-received (2000\\x02 magnification) Figure 1. SEM images of the and (b) attrition-milled ZrB2 powders (30,000\\x02 magnification).  \\x0c', 'tered body decreased by 3.4%. Then, as the B4C content increased from 0.4 to 1.0 wt.%, the relative density increased to a second maximum. However, for B4C contents above 1.0 wt.%, the relative density decreased to the minimum value observed in this study, which was 95.3%. On the whole, the relative density was more than 98.0%, except for the 4.0 wt.% B4C content sample. One reason for the variation in relative density could be that the amounts of the sintering aids could have been reduced during sintering due to their reaction with oxides that were present on the particle surfaces, which would affect the value of the calculated theoretical density. The mechanical properties of the pressureless-sintered ZrB2-SiC ceramics are also shown in Table 2. It was observed that the flexural strength, elastic modulus and Vickers hardness all had the same trend with B4C content as relative density. As the B4C concentration increased from 0 to 0.8 wt.%, the strength, modulus and hardness all decreased, then from 0.8 to 1.0 wt.% they increased to the second maximum. As B4C content increased further to 4.0 wt.%, the property values fell to a minimum due to the low relative density. These trends were nearly identical to those observed for relative density. This was in accordance with the previously analyzed effects of porosity on mechanical properties of ceramics. The materials with higher relative density had better properties. On the other hand, the fracture toughness increased as the strength, modulus and hardness decreased when more B4C added. When less B4C was added, toughness decreased due to the decrease in strength, although the reason for this was not clear. Further analysis is needed. The sintered ZrB2-SiC ceramics with 0.4 wt.% B4C additions had an average flexural strength of 382 ± 24 MPa and an elastic modulus of 405 ± 34 GPa, which was a little lower than the other pressureless sintering results for ZrB2-SiC ceramics [16]. A reasonable explanation for the low strength may be that in this study the high sintering temperature and long time hold resulted in excessive grain growth, which was harmful for the flexural strength. The low elastic modulus, which was an intrinsic material property, may have resulted from the  residual porosity existing in the microstructure [7]. The presence of a second phase (such as B4C, C, WC), microcracking or some other factors (such as an artifact of the test method) could also have influenced the elastic modulus. The Vickers hardness and fracture toughness were 3.7 ± 0.2 MPa m1/2, 13.4 ± 0.4 GPa and respectively, which were comparable to previously reported toughness values for sintered ZrB2-SiC ceramics [16]. The microstructure of the sintered ZrB2-SiC ceramics was characterized by examining the polished cross-sections of the sintered samples. The sample with 0.4 wt.% B4C inclusion was typically taken as the example. From Figure 2, it was observed that the SiC grains (dark phase) were regularly dispersed in the ZrB2 (gray phase) matrix and appeared to be elongated. This was consistent with the other reported pressureless sintering studies. The average size of SiC grains was approximately 6-8 lm, which may be attributed to the high-temperature. EDS analysis showed that the light phase contained W, which was due to attrition milling using WC balls as the milling media. The incorporated WC may be beneficial for the densification process according to other reports on its reaction with oxides on the surface of powders [25]. Because boron and carbon are light elements, it was difficult to find them using the EDS. The other samples also had analogous microstructures as shown by EPMA. At the same time, the SEM image of crack propagations from the Vickers indentation site on the polished surface of the ZrB2-SiC ceramics (not published) showed that the fracture mode was a mixture of intragranular and intergranular, and the specimens had low residual porosity. On the corner of the indention, crack deflection and branching were observed. The cracks tend to propagate through the SiC grains (intergranularly), whereas the cracks tend to propagate between ZrB2 grains (intragranularly) and obvious crack deflection was observed. Other samples with different B4C additions showed the same characteristics. In addition, the oxidation resistance of the sintered ZrB2-SiC ceramics was investigated by measuring the mass gain and the residual strength of the specimens after oxidation treatments at different temperatures. In the present study, the samples with different B4C additions were oxidized at 1400, 1500 and 1600 °C with a 30 min hold using a heating rate of 8 °C min\\x001. Results showed that, in general, as the temperature increased the mass gain also increased. However, once the B4C content was appropriate (such as 0.8-1.0 wt.%), the mass gain at 1500 °C/30 min was than that at 1400 °C/ less  Table 1. Elemental analysis of ZrB2 powders.  ZrB2 (wt.%)  B  Zr  C  O  WC  As-received  17.08  76.44  0.078  2.47  /  Attrition-milled  /  /  0.72  2.83  4.56  The amounts of B, Zr, C, O elements were measured by a conventional  chemical analysis method. The amount of WC was determined by the aver age mass change of WC milling media before and after attrition milling.  Table 2. Density and mechanical properties of ZrB2-SiC ceramics with different B4C additions.  B4C  contents  (wt.%)  Real  density (g cm\\x003)  Theoretical  density (g cm\\x003)  Relative  density  (%)  Flexural  strength  (MPa)  Elastic  modulus  (GPa)  Vickers  hardness  (GPa)  Fracture  toughness (MPa m1/2)  0  5.77  5.67  101.7  341 ± 10  382 ± 34  13.4 ± 0.2  3.8 ± 0.4  0.2  5.65  5.66  99.9  367 ± 26  410 ± 14  13.0 ± 0.8  3.5 ± 0.6  0.4  5.54  5.64  98.2  382 ± 24  405 ± 34  13.4 ± 0.4  3.7 ± 0.2  0.6  5.46  5.63  99.0  339 ± 27  360 ± 38  12.8 ± 0.4  3.4 ± 0.1  0.8  5.56  5.62  98.9  251 ± 39  333 ± 35  12.5 ± 0.6  3.0 ± 0.1  1.0  5.60  5.61  99.8  301 ± 4  391 ± 11  13.0 ± 0.2  2.5 ± 0.3  2.0  5.47  5.54  98.8  293 ± 20  333 ± 31  12.8 ± 0.3  2.8 ± 0.3  4.0  5.17  5.42  95.3  248 ± 9  307 ± 3  11.4 ± 1.0  3.7 ± 0.4  H. Zhang et al. / Scripta Materialia 60 (2009) 559-562  561  \\x0c', '562  H. Zhang et al. / Scripta Materialia 60 (2009) 559-562  dation resistance of the ZrB2-SiC ceramics improved with B4C additions between 0.4 and 1.0 wt.%. The mass gain showed a few differences at different temperatures and the residual strength was nearly 80% of the room temperature strength before oxidation.  This work was financially supported by the National Natural Science Foundation of China.  Figure  2. SEM images of  the polished cross-sections of ZrB2-SiC  ceramics with 0.4 wt.% B4C additions.  30 min. It was found that at relatively low temperatures (such as <1500 °C) the more B4C addition, the less mass gain. In contrast, when the temperature was higher (e.g. >1600 °C), B4C contents of less than 0.4 wt.% or more than 2.0 wt.% were harmful to the oxidation resistance of the ceramics. This may indicate that the addition of the correct amount of B4C (about 0.4-2.0 wt.%) could improve the oxidation resistance of the materials; otherwise the residual flexural strength was lower than the room temperature strength before oxidation. The difference was small except for the sample with the 4.0 wt.% B4C addition. For some compositions, the residual strength was higher than the unoxidized value. The residual strength and mass gain of ZrB2-SiC ceramics with 0.4 wt.% B4C additions after oxidizing at high-temperatures are shown in Figure 3. It was observed that the strength increased by 20.9% from 306 MPa (1400 °C) to (1600 °C). One 370 MPa possible reason for this increase in strength was that glassy phases could have formed at high-temperatures and healed the surface flaws, which could increase the strength. The mass gains at the three temperatures showed few differences, indicating that the silicate glassy phase on the surface may be more effective at reducing mass gain at 1600 °C, and possibly higher temperatures. As a whole, the addition of the correct amount of B4C improved the oxidation resistance of the materials to some extent whereas excessive B4C addition deteriorated the oxidation resistance due to the evaporation of B2O3. In conclusion, ZrB2-SiC ceramics were densified to (\\x18100%) nearly full density by pressureless sintering with different B4C additions. It was shown that as the addition of B4C increased, the relative density of the sintered bodies first decreased, then increased and finally fell to a minimum. The mechanical properties, such as the flexural strength and elastic modulus, followed the same trend as the relative density. In addition, the oxi Figure 3. Dependence of mass gain and flexural strength of ZrB2-SiC  ceramics with 0.4 wt.% B4C additions 1600 °C for 30 min.  oxidized  at  1400,  1500  or  et  et  al.,  J. Am.  Process.  al.,  J. Am.  Ion 172  (1-4)  J. Mater.  [1] R.A. Cutler, in: S.J. Schneider (Ed.), Ceramics and Glasses, Engineered Materials Handbook, vol. 4, ASM International, Materials Park, OH, 1992, pp. 787-803. [2] C. Mroz, Am. Ceram. Soc. Bull. 73 (6) (1994) 141-142. [3] R. Telle, L.S. Sigl, et al., in: R. Riedel (Ed.), Handbook of Ceramic Hard Materials, Wiley-VCH, Weinheim, 2000, pp. 802-945. [4] S.R. Levine, E.J. Opila, et al., J. Eur. Ceram. Soc. 22 (1415) (2002) 2757-2767. [5] F. Monteverde, A. Bellosi, et al., J. Eur. Ceram. Soc. 22 (3) (2002) 279-288. [6] A.L. Chamberlain, W.G. Fahrenholtz, Ceram. Soc. 87 (6) (2004) 1170-1172. [7] A.L. Chamberlain, W.G. Fahrenholtz, Ceram. Soc. 89 (2) (2006) 450-456. [8] K. Upadhya, J.M. Yang, et al., Am. Ceram. Soc. Bull. 76 (12) (1997) 51-56. [9] F. Monteverde, S. Guicciardi, et al., Mater. Sci. Eng. A - Struct. Mater. Prop. Microstruct. Process. 346 (1-2) (2003) 310-319. [10] S.N. Karlsdottir, J.W. Halloran, J. Am. Ceram. Soc. 90 (10) (2007) 3233-3238. [11] A. Bellosi, F. Monteverde, et al., Manuf. Sci. 9 (2) (2000) 156-170. [12] T. Tsuchida, S. Yamamoto, Solid State (2004) 215-216. [13] Y. Zhao, L.J. Wang, et al., J. Am. Ceram. Soc. 90 (12) (2007) 4040-4042. [14] G.J. Zhang, Z.Y. Deng, et al., J. Am. Ceram. Soc. 83 (9) (2000) 2330-2332. [15] J.W. Zimmermann, G.E. Hilmas, Soc. 27 (7) (2007) 2729-2736. [16] C. Zhang Shi, E. Hilmas Greg, et al., J. Am. Ceram. Soc. 91 (1) (2008) 26-32. [17] F. Monteverde, A. Bellosi, Scr. Mater. 46 (3) (2002) 223- 228. [18] F. Monteverde, A. Bellosi, Solid State Sci. 7 (5) (2005) 622- 630. [19] A. Rezaie, W.G. Fahrenholtz, et al., J. Mater. Sci. 42 (8) (2007) 2735-2744. [20] Y.J. Yan, Z.R. Huang, et al., J. Am. Ceram. Soc. 89 (11) (2006) 3589-3592. [21] S.K. Mishra, S.K. Das, et al., J. Am. Ceram. Soc. 85 (11) (2002) 2846-2848. [22] J.J. Melendez-Martinez, A. Dominguez-Rodriguez, et al., J. Eur. Ceram. Soc. 22 (14-15) (2002) 2543-2549. [23] S.C. Zhang, G.E. Hilmas, et al., J. Am. Ceram. Soc. 89 (5) (2006) 1544-1550. [24] S.M. Zhu, W.G. Fahrenholtz, et al., Mater. Sci. Eng. A - Struct. Mater. Prop. Microstruct. Process. 459 (1-2) (2007) 167-171. [25] S. Zhu, W.G. Fahrenholtz, et al., J. Am. Ceram. Soc. 90 (11) (2007) 3660-3663. [26] W.G. Fahrenholtz, J. Am. Ceram. Soc. 90 (1) (2007) 143- 148. [27] D. Sciti, M. Brach, et al., J. Mater. Res. 20 (4) (2005) 922- 930. [28] D. Sciti, S. Guicciardi, et al., J. Am. Ceram. Soc. 89 (7) (2006) 2320-2322.  J. Eur. Ceram.  et al.,  \\x0c']"
},{
  "_id": 217,
  "PDF": "Pressureless sintering, mechanical properties and oxidation behavior of ZrB2 ceramics doped with B4C.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 35 (2015) 2699-2705  Pressureless sintering, mechanical properties and oxidation behavior of ZrB2 ceramics doped with B4C  Hai-Bin Ma, Hu-Lin Liu, Jian Zhao, Fang-Fang Xu, Guo-Jun Zhang  ∗  State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Shanghai 200050, China  Received 19 January 2015; received in revised form 16 March 2015; accepted 21 March 2015  Available online 9 April 2015  Abstract     A pressureless sintering process using different amounts of boron carbide (B4C) as sintering aid was investigated for the densification of zirconium diboride (ZrB2 ). The sintering mechanism was converted from solid state sintering to liquid phase sintering as sintering temperature increased from 2150 to 2250 C. B4C addition improved densification by removing the oxide impurities during solid state sintering. Besides, it could restrain the growth of ZrB2 grains. At 2250 C, ZrB2 and B4C formed liquid phase which was well wetting ZrB2 grains during liquid phase sintering process. The different roles of B4C on the microstructure development, mechanical properties and oxidation resistance of the sintered ZrB2 -based ceramics were investigated. © 2015 Elsevier Ltd. All rights reserved.     Keywords: ZrB2 ; Pressureless sintering; Microstructure; Mechanical property; Oxidation behavior  1.   Introduction  Zirconium diboride  (ZrB2 ) and Hafnium diboride  (HfB2 ) were especially promising materials for high temperature structural  applications,  such  as  thermal protection  structures  for hypersonic vehicles, molten metal crucibles, cutting  tools and even heating elements, owing to the combination of high melting temperature, high  strength, high  thermal  and  electrical  conductivity  [1]. Different processes  such  as hot pressing  (HP) [2-4],  reactive hot pressing  (RHP)  [5-7]  and  spark plasma sintering  (SPS)  [8-11] have been used  to densify ZrB2 -based ultra-high  temperature  ceramics  (UHTCs). Compared with pressure-assisted sintering, pressureless sintering  is more convenient and allows  fabricating components  to near-net shape. However, pressureless  sintering of  transition metal diborides has been proved very difficult. Therefore, some sintering aids are usually chosen to help the densification process of UHTCs. Sintering aids used in recent years can be classified into three kinds. One kind  is  transition metals such as Fe, Cr  [12] and Ni [13], which can form liquid phases during sintering process.     Then,  these  liquid phases can fill voids among matrix grains and favor  the mass  transport  that can accelerate  the densification process. Another is refractory metal silicides such as MoSi2 [14], TiSi2 [15] and HfSi2 [16]. Sciti et al. [14] successfully prepared dense HfB2 -based ceramics using 5-20 vol.% MoSi2 as sintering aids at 1900-1950 C via pressureless sintering. They thought that MoSi2 was softened at high temperature which had the similar effect with transition metals. The third kind, which is also used intensively, is reductants such as C [17,18], B [19], B4C [17,20], WC [21,22], VC [23] and their mixtures [17,24]. These aids can react with oxide  impurities and clean  the particle surface. Therefore, nearly fully dense ZrB2 -based and HfB2 -based ceramics can be prepared without external pressure. Among the numerous sintering aids, WC and B4C are more suitable to be selected. For WC, it has excellent oxide removing ability  [25]. Furthermore,  it has been  reported  that  the addition of ‘W’ can improve the oxidation resistance of ZrB2 -based ceramics  [1]. For B4C,  taking ZrB2 as an example,  the  reaction product between ZrO2 and B4C  is ZrB2 itself, as shown in  reaction  (1). This might  improve mechanical properties of ZrB2 -based ceramics at elevated temperature [17,20].  ∗  Corresponding author. Tel.: +86 21 52411080; fax: +86 21 52413122.  E-mail address: gjzhang@mail.sic.ac.cn (G.-J. Zhang).  7ZrO2 +   5B4C   =   7ZrB2 +   B2O3 (l,   g)   +   5CO(g)   (1)  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.03.030  0955-2219/© 2015 Elsevier Ltd. All rights reserved.        \\x0c', '2700   H.-B. Ma et al. / Journal of the European Ceramic Society 35 (2015) 2699-2705  Fig. 1. Relative densities (a) and closed porosities (b) of ZrB2 with different content of B4C pressureless sintered at 1950-2300     C.  A number of investigations have explored the effect of B4C on  the  sintering behavior and mechanical properties of ZrB2 and HfB2 . However, most of them on the densification of ZrB2 and HfB2 were  conducted below 2200 C  and believed  that solid state sintering was  the main sintering mechanism when using B4C  as  sintering  aid. Zou  et  al.,  [26]  found  that  the sintering mechanism  converted  from  solid  state  sintering  to liquid phase  sintering when  temperature was above 2300 C in HfB2-B4C  system. Similarly, ZrB2-B4C  system also has a eutectic  temperature at 2220   20 C according  to Zr-C-B phase diagram  [27]. Therefore,  the  role of B4C  at different  ±           temperatures might be different  in ZrB2 -B4C  system.  In  the present work, ZrB2-B4C ceramics were  sintered at different sintering  temperatures. Then,  the different roles of B4C on  the sintering behavior and microstructure of ZrB2 were studied. The mechanical properties and oxidation resistance of the as-sintered ZrB2 -based ceramics were also characterized and discussed.  2. Experimental procedure  The  raw materials  used were ZrB2 powder  synthesized [28]  through a boron/carbon  thermal  reaction between ZrO2  Fig. 2. Microstructures of ZB5 pressureless sintered at (a) 1950     C, (b) 2050     C, (c) 2150     C, (d) 2250     C, (e-g) 2300     C, (i) ZB0 at 2300     C, (h) EDS pattern of the  black phase in (g).  \\x0c', 'H.-B. Ma et al. / Journal of the European Ceramic Society 35 (2015) 2699-2705   2701          ×  ×  and B4C  in vacuum  (D50 = 1.05  \\u242em, purity 98%, O 0.46%, C 0.1%, Hf 1.10%, others < 0.1%) and commercial B4C powder (D50 = 1.5  \\u242em, Jingangzuan Boron Carbide Co. Ltd., Mudanjiang, China). The content of B4C was 0, 1, 3, and 5 vol.% and  the corresponding samples were designated as ZB0, ZB1, ZB3, and ZB5 for simplicity. The starting mixture was mixed for 24 h  in a polyethylene  jar using ethyl alcohol and Si3N4 balls. Then  the powders were dried by rotary evaporation and sieved through a 200-mesh screen. Rectangle pellets with dimensions of 37   30   5 mm were shaped by uniaxial pressing at 30 MPa, followed by cold isostatic pressing (CIP) at 200 MPa for 2 min. Sintering was carried out in a high temperature graphite resistant furnace (MRF3338, Materials Research Furnaces Inc., Suncook, USA) at temperatures ranging from 1950 to 2300 C for 2 h. The compacts were sintered  in vacuum  from  room  temperature  to 1600 C, and continually heated to 1950, 2050, 2150, 2250 and 2300 C in high purity argon atmosphere. The  relative density of  the sintered samples was measured by  the Archimedes method  in  deionized water. Four-point bending  strength was measured on bars with dimensions of 2   2.5   30 mm. One  tensile  surface of each  specimen was polished and  four edges were chamfered. The bars  for oxidation  testing were cut from  the sintered billets with dimensions of 2   2.5   15 mm. Samples were placed on a zirconia support with  ridges  to minimize  the contact area. And oxidation  tests were conducted in a muffle furnace (1708 BL, CM Furnaces Inc., USA). Samples were heated  to 1500 C with a heating rate of 5 C/min and the dwelling time was 0 h in stagnant air. Then the furnace was cooled naturally. Mass of the samples was measured using a balance with 0.1 mg precision before and after oxidation. Microstructures and chemical composition were analyzed using scanning electron microscopy  (SEM, TM3000, Hitachi, Japan) along with energy-dispersive spectroscopy (EDS, Oxford INCA energy). Phase composition was  investigated via  transmission electron microscopy (TEM, Tecnai-F20, FEI Co., USA) on as-sintered samples.  ×  ×  ×  ×        3. Results and discussion  3.1. Densification behavior     The relative densities of pure ZrB2 and ZrB2 with different amounts of B4C sintered  in  temperature  range 1950-2300 C are shown in Fig. 1(a). As the calculation of theoretical densities was based on the initial composition, the relative densities shown here were not precise. However, it could still reflect the varying trend of densities.  It was  found  that  the densification of pure ZrB2 was difficult just by increasing sintering temperature. The 85%. highest relative density of ZrB2 without B4C was only  The addition of B4C could accelerate the densification of ZrB2 ceramics. The ZB5 sintered at 2300 C had the highest relative density, 97.12%. The closed porosities of ZB1, ZB3 and ZB5 are shown  in Fig. 1(b).  It  could be  found  that  the  closed porosities were in accordance with  the relative densities. However,  2% which the closed porosity of ZB1  sintered  at 1950 C was only  was much  lower  than others. Because of  the  low  sintering        temperature and low content of B4C for ZB1, the porosity was very high and  resulted  in a pore connection with a high open porosity of 22.53%. Therefore, its closed porosity was very low. The relative densities and closed porosities of samples sintered at 2150 C were nearly equal  to  those of samples sintered at 2250 C. However, the locations of pores in bulk specimens were different.        3.2. Microstructure           ZrB2 doped with 5 vol.% B4C demonstrated  the best  sinterability, so here  its microstructure evolution combined with different sintering  temperatures was focused  to  investigate and the  results  are  shown  in  Fig.  2. As  sintering  temperature increased from 1950 to 2150 C, the quantity and size of pores decreased which was  in accordance with  the value of  relative density. Pores still existed as sintering  temperature  increasing to 2250 and 2300 C. However,  the position of pores  in ZB5 at 2150 C is obviously different from that in ZB5 at 2250 and 2300 C. Most of pores in ZB5 sintered at 2150 C were in the triple  junctions and  few pores  in  the grains. But  for ZB5 sintered 2250 and 2300 C, almost all the pores are in the grains, as shown with white arrows in Fig. 2(f) and (g). Besides, some B4C particles were entrapped  into  these grains  (shown with black arrows). The grain size distributions of ZrB2 phase sintered at different temperatures are shown  in Fig. 3 and  the average grain sizes of  them are  listed  in Table 1. Mild grain growth of ZrB2 with           Fig. 3. Grain size distributions of ZB5 sintered at different   temperatures:   (a)                 1950  C, (b) 2050  C, (c) 2150  C, (d) 2250  C and (e) 2300  C.  \\x0c', '2702   H.-B. Ma et al. / Journal of the European Ceramic Society 35 (2015) 2699-2705  Table 1  Average grain size of ZrB2 with different amounts of B4C at different sintering temperatures.  Amount of B4C (vol.%)  Average grain size of ZrB2 (\\u242em)  2300     C  2250     C   2150     C   2050     C   1950     C  0   3.68   ± ± ± ±   1.32   -   -   -   -  1   46.13    21.60  40.33   ± ± ±   18.90  21.76   ± ± ±   12.01  6.63   ± ± ±   2.50  2.73   ± ± ±   0.97  3   77.88    36.24  62.07    23.93  20.46    10.39   12.01    6.16   6.07    2.67  5   94.90    34.25   70.13    29.15   18.53    8.16   7.98    3.45   4.94    1.99  Table 2  Average grain size of B4C at different sintering temperatures.  Amount of B4C (vol.%)   Average grain size of B4C (\\u242em)  1950     C   2050     C   2150     C  1   1.53   1.81   3.63  3   1.71   2.00   4.01  5   1.70   1.85   3.47  narrow grain size distribution was observed  in  the  temperature range of 1950-2050 C. The grains grew a little larger at 2150 C with a similar morphology as  that at 1950-2050 C. However, grain growth appeared  to be rapid at 2250 C, and  the average \\u242em, which was much  grain size was 70.13   29.15  larger  than that at 2150 C. This was caused by  the sintering mechanism from solid state sintering to liquid phase sintering, which will be discussed in detail in next paragraph. Besides, the morphology of ZrB2 at 2250 C was quasi-spherical, different from that of ZrB2 at 1950-2150 C which was equiaxial. When  the  temperature increased  to 2300 C,  the grain size of ZrB2 increased further and the morphology kept quasi-spherical.              ±              3.3. Sintering mechanism  For ZrB2 , oxide impurities are present on the particle surfaces of the starting powders and reduce the surface energy of powders, which have adverse effect on the sintering properties. As a result, pure ZrB2 ceramics without B4C addition had  residual pores in spite of  the sintering  temperature  increasing  to 2300 C, as shown  in Fig. 2(h). The driving  force  for  sintering could be evaluated by  the ratio of  the surface energy (γ SV ) of a powder compact and the grain boundary energy (γ GB ). Higher  indicates  the  larger driving  force and better sinterability. The added B4C could react with ZrO2 at high temperature, and      γ SV /γ GB  γ SV  increased due  to  the purified surfaces of ZrB2 powders. More oxygen contamination would be eliminated with the increasing B4C amount, due to the greater contact probability between B4C and ZrO2 on the surfaces of ZrB2 powders. When sintering  temperature was at  the range from 1950 C to 2150 C,  the sintering mechanism was solid state sintering. As mentioned before, a  relatively mild grain growth behavior was observed  for  the process of solid state sintering. At  this temperature range, B4C has  two roles for  the sintering process of ZrB2-xB4C. Besides being a reductant reacting with oxygen contamination, it could restrain grain growth as a second phase. When sintering temperature was below 2150 C, the B4C particles were isolated and uniformly dispersed in the matrix of ZrB2 . According  to Smith and Zener-based models [29],  the  limited grain size of major phase was defined as Eq. (2).           D = k  × d  f  (2)  where D, d and f represent the average diameter of matrix grains, the average diameter and the volume fraction of the added second phase particles, respectively. According to Table 2, average grain size of B4C did not change significantly as  the amounts increased from 1 vol.% to 5 vol.%. Therefore, the average grain size of ZrB2 should have decreased when  amounts of B4C increased. However,  in  fact,  the  average grain  size of ZrB2 increased obviously as B4C increased from 1 to 3 vol.% at 1950 and 2050 C. When the amount of B4C was only 1 vol.%, oxygen contaminations could not be removed due to the limited contact between B4C and ZrO2 on the surfaces of ZrB2 powders. Therefore, at  low sintering  temperature,  the diffusion of ZrB2 was very difficult leading to small average grain size of ZrB2 . When B4C amounts were above 3 vol.%, most oxygen contamination was eliminated, and the residual B4C could pin the ZrB2 grain     Fig. 4.   (a) TEM image of ZB5 sintered at 2250     C and (b) corresponding ED pattern of B4C.    \\x0c', 'H.-B. Ma et al. / Journal of the European Ceramic Society 35 (2015) 2699-2705   2703        ±  boundaries. Under this condition, the average grain size of ZrB2 was in accordance with Smith and Zener-based models. ZrB2-B4C  system  has  the  eutectic  temperature  at 2220   20 C  according  to  Zr-C-B  phase  diagram  [27]. When sintering temperature increased to 2250 C, liquid phase was  formed by  the  reaction  (3) between ZrB2 and B4C.  In equilibrium state, all  the B4C was consumed and converted  to the liquid phase when the reaction (4) completed. Such kind of liquid phase is effective on enhancing the densification process, because  the  surface of ZrB2 grains was  totally wetted by  it and  the mass  transport was accelerated. EDS results  indicated that  the black phase at  the  triple  joints contained B, C and Zr  elements  (shown  in Fig. 2(i)), which was  in  accordance with Zr-C-B phase diagram. It also could be found from EDS results  that  the amount of Zr element was very  low (less  than 1 at.%). Therefore, we  speculated  that  the black phase was B4C. The amount of black phase was so  low  that only ZrB2 phase could be detected via X-ray diffraction (XRD). TEM was used  to  identify  the state of  the black phase. TEM  images and electronic diffraction  (ED) pattern were shown  in Fig. 4. The results of ED indicated that the black phase in the triple points was single crystalline B4C. ZrB2 +  ZrB2 + ZrB2 +  + (cid:5) +   ZrB2  = =   B4C    B4C    B4C   (3)  (4)   L    L  ␥GB and   When sintering temperature increased from 2150 to 2250 C, the  liquid  phase was  formed  and wetted  the ZrB2 grains. This  resulted  in  the  reduction of  increase of  sintering driving  force. Accordingly,  the  sintering mechanism was converted  from  solid  state  sintering  to  liquid  phase  sintering. When  the  sintering  temperature was higher,  the  liquid phase  could better wet  the ZrB2 grains. This was  in  accordance with  the value of dihedral angle between ZrB2 grains and  the  liquid phase at different  temperatures. The measured average dihedral angle between ZrB2 grains and the liquid phase was 15.6   6.7 when sintered at 2300 C, slightly  lower  than that sintered at 2250 C  (19.7   9.3 ). On  the other hand,  the formed  liquid can also accelerates  the mass  transfer via  the dissolution-diffusion-precipitation process. The smaller grains were dissolved and precipitated on  the  larger grains. The process resulted in the increase of grains size and the quasi-spherical morphology as shown in Table 1 and Fig. 2(d) and (e).  ±  ±                 3.4. Properties  3.4.1. Mechanical properties  Bending  strength of  some  selected ZrB2-xB4C  ceramics were summarized  in Fig. 5. According  to Hall-Petch relationship (5):  σy =  σ0 +  Kd−1/2  (5)  where  σ 0 is a certain shear stress required to ensure gliding dislocations in a monocrystal, and K is a constant specific to each material also known as the Hall-Petch parameter. As discussed above,  liquid phase  sintering  leading  to  rapid grain growth. Therefore, the bending strength of ZrB2 -xB4C sintered at 2250  Fig. 5. Bending strength of ZrB2 -B4C pressureless sintered at 1950-2300 C. −1/2 of (b) (a) bending strength versus B4C amount; bending strength versus d ZrB2 and (c) B4C.           and 2300 C (lower than 300 MPa) was lower than that sintered at 1950-2150 C. As  the content of B4C  increased  from 1  to 5 vol.%,  the bending strengths of samples sintered at 2250 C decreased  from  273   36 MPa  to  157   24 MPa  because  of increasing grain size. And samples sintered at 2300 C had the same  tendency as  samples  sintered at 2250 C. The  relation −1/2 ) of ZrB2 and bending strength between  the grain size  (d −1/2 and (σ ) was shown  in Fig. 5(b).  It could be  found  that d  ±  ±               \\x0c', '2704   H.-B. Ma et al. / Journal of the European Ceramic Society 35 (2015) 2699-2705  Fig. 6. Surface morphology of the oxidized (a) ZB5-2150 and (b) ZB5-2250.  σ     nearly  satisfied  linear  relation. Which  indicated  that  the relation between bending strength and grain size of ZrB2 for ZrB2-B4C  sintered at 1950-2300 C was  in  full accordance with Hall-Petch  relationship. Previous  results  [1] showed  the SiC grain size, instead of that in ZrB2 , played a determinant role on strength  in ZrB2-SiC system. In present work,  the relation −1/2 ) of B4C and bending strength (σ ) between the grain size (d was also shown (Fig. 5(c)). From the results, in ZrB2-B4C system in this work, the B4C grain size did not play a determinant role on strength.  3.4.2. Oxidation properties              Oxidation resistance is a very important property for UHTCs. Here, we  studied  the oxidation  resistance of ZB5  sintered at 2150 and 2250 C (designated as ZB5-2150 and ZB5-2250 for simplicity). The  relative densities of  these  two samples were nearly equal (95.94% for ZB5-2150 and 95.77% for ZB5-2250). The grain size and morphology of ZrB2 and B4C were obviously different  from each other. Therefore,  these  two samples were chosen to study the effect of microstructural morphology on the oxidation resistance. 800 It has been  reported  [30]  that ZrB2 begins  to oxidize  at C. The  time was 140 min  from 800  to 1500 C during heating stage and  the  time was about 30 min when  the furnace cooled to 800 C. Therefore, the total oxidation time was about 3 h. Mass gain per unit  surface area of ZB5-2150 and ZB52250 was 0.05 and 0.11 mg/cm2 , respectively. The thickness of 193 and  287  oxidation  layer was  \\u242em. These  results  indicated that ZB5-2150 had better property of oxidation resistance than ZB5-2250. Surface morphologies of  the oxidized specimens are shown in Fig. 6. As above-discussed, the morphology of B4C changed obviously when sintering temperature increased from 2150 to 2250 C. During the oxidation process, ZrB2 and B4C react with oxygen (O2 ), as show  in reactions (6) and (7). B2O3 volatilizes when  temperature approaches 1200 C  [31]. For ZB5-2150, B4C particles were small and isolated, resulting in a relatively dense oxidation layer of ZrO2 (shown in Fig. 6(a)). However,  for ZB5-2250, most B4C were at  triple points and some of  them were connected  through  the grain boundaries. Therefore, large holes (shown in Fig. 6(b)) generated at the triple points during oxidation process which acted as O2 diffusion channels, resulting in a faster oxidation.        O2 (g)   =   ZrO2 +   B2O3 (l,   g)   ZrB2 + 5 2  +  B4C    4O2 (g)   =   2B2O3 (l,   g)   +   CO2 (g)   (6)  (7)  4. Conclusions        Zirconium diboride ceramics were densified by pressureless sintering with B4C as additive (1, 3, 5 vol.%) at 1950-2300 C for 2 h. When  sintering  temperature  increased  from 2150  to 2250 C, the sintering mechanism was converted from solid state sintering to liquid phase assisted sintering. For solid state sintering, B4C acted as a second phase to restrain ZrB2 grains growth. However, for liquid phase sintering, increasing amount of B4C generated more  liquid  leading  to faster grain growth. As grain growth was  rapid at 2250 and 2300 C, some pores and B4C particles were entrapped  in  the grains of ZrB2 . Liquid phase was beneficial to the densification process, whereas lager grain size of ZrB2 was detrimental  to mechanical properties. During oxidation at 1500 C, the gathered B4C phase at the triple points generated large holes which acted as oxygen diffusion channels and resulting  in a faster oxidation process. ZrB2 -5 vol.% B4C sintered at 2050 C owns  the best flexural strength (470 MPa) and oxidation resistance.           References  [1]. Fahrenholtz WG, Hilmas GE, Talmy IG, Zaykoski JA. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90:1347-64.  [2]. Sciti D, Monteverde F, Guicciardi S, Pezzotti G, Bellosi A. Microstructure  and mechanical properties of ZrB2 -MoSi2 ceramic composites produced by different sintering   techniques. Mater Sci Eng A - Struct Mater Prop Microstruct Process 2006;434:303-9.  [3]. Zou J, Zhang GJ, Kan YM, Wang PL. Hot-pressed ZrB2 -SiC ceramics with VC addition: chemical reactions, microstructures, and mechanical properties. J Am Ceram Soc 2009;92:2838-46.  [4]. Guo WM, Zhou XJ, Zhang GJ, Kan YM, Li YG, Wang PL. Effect of Si and  Zr additions on oxidation resistance of hot-pressed ZrB2 -SiC composites with polycarbosilane as a precursor at 1500 degrees C. J Alloys Compd 2009;471:153-6.  [5]. Zhang GJ, Deng ZY, Kondo N, Yang JF, Ohji T. Reactive hot pressing of ZrB2 -SiC composites. J Am Ceram Soc 2000;83:2330-2. [6]. Wu WW, Zhang GJ, Kan YM, Wang PL, Vanmeensel K, Vleugels   J,  et al. Synthesis and microstructural   features of ZrB2 -SiC-based composites by reactive spark plasma sintering and reactive hot pressing. Scr Mater 2007;57:317-20.    \\x0c', 'H.-B. 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Effect   of Fe   and  Cr   addition   on   the   sintering   behavior   of   ZrB2  produced   by  self2002;85:  propagating   high-temperature   synthesis.   J   Am   Ceram   Soc   2846-8.  [13].Monteverde   F,   Bellosi A, Guicciardi   S.   Processing   and   properties 2002;22:  of   zirconium   diboride-based   composites.   J Eur Ceram   Soc   279-88.  [14]. Silvestroni L, Sciti D. Effects of MoSi2 Hfand Zr-B2 composites produced by pressureless sintering. Scr Mater 2007;57:165-8.  additions on   the properties of  [15]. Sonber JK, Murthy TSRC, Subramanian C, Kumar S, Fotedar RK, Suri AK.  Investigations on synthesis of HfB2 and development of a new composite  with TiSi2 . Int J Refract Met Hard Mater 2010;28:201-10.  [16].Monteverde F. Hot pressing of hafnium diboride aided by different sinter additives. J Mater Sci 2008;43:1002-7.  [17]. Fahrenholtz WG, Hilmas GE, Zhang SC, Zhu S. Pressureless sintering of  zirconium diboride: particle  2008;91:1398-404.  size and additive effects. J Am Ceram Soc  [18]. Zhu S, Fahrenholtz WG, Hilmas GE, Zhang SC. Pressureless sintering of  carbon-coated zirconium diboride powders. Mater Sci Eng A - Struct Mater  Prop Microstruct Process 2007;459:167-71.  [19]. Guo WM, Zhang GJ, Yang ZG. Pressureless sintering of zirconium diboride ceramics with boron additive. J Am Ceram Soc 2012;95:2470-3.  [20]. Zhang SC, Hilmas GE, Fahrenholtz WG. Pressureless densification of  zirconium  diboride with  2006;89:1544-50.  boron   carbide   additions.   J   Am Ceram   Soc  [21]. Chamberlain AL, Fahrenholtz WG, Hilmas GE. Pressureless sintering of zirconium diboride. J Am Ceram Soc 2006;89:450-6.  [22]. Zou J, Zhang GJ, Kan YM. Formation of tough interlocking microstructure  in ZrB2 -SiC-based ultrahigh-temperature ceramics by pressureless sintering. J Mater Res 2009;24:2428-34.  [23]. Zou J, Zhang GJ, Kan YM, Wang PL. Pressureless densification of ZrB(2)SiC composites with vanadium carbide. Scr Mater 2008;59:309-12.  [24]. Zhu S, Fahrenholtz WG, Hilmas GE, Zhang SC. Pressureless sintering of  zirconium diboride using boron carbide and carbon additions. J Am Ceram Soc 2007;90:3660-3.  [25]. Zou J, Zhang GJ, Sun SK, Liu HT, Kan YM, Liu JX, et al. ZrO2 ing  reactions of Groups  IV-VI  transition metal carbides  in ZrB2 based composites. J Eur Ceram Soc 2011;31:421-7.  remov [26]. Zou J, Zhang GJ, Kan YM. Pressureless densification and mechanical prop erties of hafnium diboride doped with B4C:  from solid state sintering  liquid phase sintering. J Eur Ceram Soc 2010;30:2699-705.  to  [27]. Rudy E, Windisch S. Part II. Ternary systems. Volume XIII: Phase diagrams  of   the systems Ti-B-C, Zr-B-C and Hf-B-C. Ternary phase equilibria   in  transition metal-boron-carbon-silicon systems. Report No. AFMLTR-65 2. Contract No. USAF 33(615)-1249. Ohio: Air Force Materials Laboratory,  Wright-Patterson Air Force Base; 1966. p. 183-8.  [28]. Zou J, Zhang GJ, Vleugels J, Van der Biest O. High   temperature strength  of hot pressed ZrB2 -20 vol% SiC ceramics based on ZrB2 starting powders prepared by different carbo/boro-thermal reduction routes. J Eur Ceram Soc 2013;33:1609-14.  [29]. Rohrer GS. Introduction to grains, phases, and interfaces—an interpretation  of microstructure, Trans. AIME, 1948, vol. 175, pp. 15-51, by C.S. Smith.  Metall Mater Trans A 2010;41:1063-100.  [30]. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during   the  oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27:2495-501.     C. J Eur  [31].Monteverde F, Bellosi A. Oxidation of ZrB2 -based ceramics  Electrochem Soc 2003;150:B552-9.  in dry air. J  \\x0c']"
},{
  "_id": 218,
  "PDF": "Processing and properties of ZrC, ZrN and ZrCN ceramics- a review.pdf",
  "Text": "['Advances in Applied Ceramics  Structural, Functional and Bioceramics  ISSN: 1743-6753 (Print) 1743-6761 (Online) Journal homepage: https://www.tandfonline.com/loi/yaac20  Processing and properties of ZrC, ZrN and ZrCN ceramics: a review  R. W. Harrison & W. E. Lee  To cite this article: R. W. Harrison & W. E. Lee (2016) Processing and properties of ZrC, ZrN and ZrCN ceramics: a review, Advances in Applied Ceramics, 115:5, 294-307, DOI: 10.1179/1743676115Y.0000000061  To link to this article:  https://doi.org/10.1179/1743676115Y.0000000061  Published online: 06 May 2016.  Submit your article to this journal   Article views: 5262  View related articles   View Crossmark data  Citing articles: 37 View citing articles   Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=yaac20  \\x0c', 'Processing and properties of ZrC, ZrN and  ZrCN ceramics: a review  R. W. Harrison*1,2 and W. E. Lee1  ZrC and ZrN ceramics are of  interest  in the application of materials in extreme high temperature  environments, particularly for nuclear applications in generation IV reactors. These materials  demonstrate desirable characteristics such as high thermal and electrical conductivities along  with high hardness and melting temperatures. Data reported in the literature often suffer  from  scatter due to differences in processing techniques and difficulty determining stoichiometry,  which will significantly affect  thermophysical properties. This article reviews the current available  data for  the properties of ZrC, ZrN and mixed carbonitrides phases and identifies causes of  scatter in the literature and areas requiring further research.  Keywords: Zirconium carbide, Zirconium nitride, Properties, Review, Processing  Introduction  Non-oxide ceramics such as the carbides and nitrides of  group IV transition metal elements have an interesting  mixture of  ionic,  covalent and metallic bonding. This  combination of bonding gives the materials an unusual  mixture of advantageous properties such as high hard(*25 GPa) ness and very high melting temperatures (>3000 K) along with good thermal and electrical conductivity (>10 W m and *200|104 V21 m respectively).1 As a result of these properties, zirconium  21 K 21  21  carbide and nitride are receiving increased attention for  use in advanced nuclear power plants (NPPs) and as de  facto ultrahigh temperature  ceramics  for other  appli cations in extreme environments. However,  the proces sing  routes  to  these materials  can  lead  to  varying  contents of impurities such as oxygen in ZrC and oxygen  and  carbon  in ZrN. Both materials  can  also  easily  accommodate vacancies on the carbon or nitrogen site in  the lattice. The presence of these point defects inevitably  affects  the  thermal  and mechanical properties of  the  materials and can be a source of scatter in the literature.  This article collates and presents data from the 1960s up  to present  efforts  in the  research fields of  fabrication,  processing and characterisation of ZrC, ZrN and ZrCN  ceramics.  Non-oxide processing  Powder fabrication  Synthesis  of  carbide  and  nitride materials  generally  involves direct reaction with the metal, the metal hydride  or the metal oxide, with the metal oxide route being the  main  precursor  from  spent  fuel  reprocessing.  Other  techniques have been examined including propagating high temperature synthesis2,3 chemical vapour deposition techniques.4-6  self as well  as  Methods of carbothermal reduction and nitridation of processing parameters,7,8  oxides  generally use  similar  mixing  of  oxide  powder with  a  carbon  source,  heat  treatment  to between 1873 and 2373 K under  inert at mosphere, and then further heat treatment in a hydrogen  doped nitrogen atmosphere to complete the reaction and  remove  excess  carbon. The  basic mechanism of  car bothermic  reduction  and  nitridation  is  as  follows  (equations (1) and (2)): MO2 þ 3C ! MC þ 2CO  ð1Þ  MC þ 1 2  N2 þ 2H2 ! MN þ CH4  ð2Þ  The mechanisms of carbothermal reduction of ZrO2 and  TiO2, which are a similar system to U/PuO2, have been mechanism for ZrO2 þ xC begins with CO formation from examined by Berger et al.9 The authors propose that the the solid state reactions of oxide and carbon particles. The  increased number of  lattice defects arising from their for mation in the substoichiometric ZrO2-x allows incorporation of carbon into the vacancies, leading to a ZrCxOy phase. The  authors state that the reaction rates of oxygen loss and car bide formation are not equal, which leads to substoichio metric oxycarbides. The substitution of oxygen for carbon  proceeds via the same mechanism of oxygen removal of O by  CO to give CO2 and the incorporation of C into the vacancies left behind. This is a similar mechanism to that reported by Pautasso et al.10 and Mukerjee et al.11 whereby C and N  occupy vacancies left from oxygen removal.  TEM studies of  the carbothermal reduction of zirco nia to assess the reaction mechanisms have been reporal.12  ted  by  David  et  Samples  of  ZrO2 ratio of 2.89 were  and  C  (carbon black) with a C/Zr molar  heated at temperatures of 1873 and 2023 K for dwells of  1Centre for Nuclear Engineering and Department of Materials,  Imperial  College London, London SW7 2AZ, UK 2Department of Computing and Engineering, University of Huddersfield,  Queensgate, Huddersfield HD1 3DH, UK  *Corresponding author, email R.W.Harrison@Hud.ac.uk  Ñ 2 0 1 6 I n s t i t u t e o f M a t e r i a l s , M i n e r a l s a n d M i n i n g  R e c e i v e d 1 4 M a y 2 0 1 5 ; a c c e p t e d 13 J u l y 2 0 1 5  DO I 1 0 . 1 1 7 9 / 17 4 3 6 7 6 1 1 5 Y . 0 0 0 0 0 0 00 6 1  294  Pub l ished by Tay lo r & F ranc is on beha l f o f  the Ins t i tu te  REVIE s  sW  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  s s s s \\x0c', '0-8 h. TEM studies on powders  that had undergone  minimal  reaction showed that ZrCxOy  is not observed  on ZrO2 particles but located in the vicinity of them. The  authors state that  the mechanism begins with destabili sation of carbon black and zirconia, with oxidation of  carbon leading to CO(g) and the removal of oxygen from  zirconia giving a ZrO2-x(s) phase. This destabilisation of  zirconia continues until a layer of Zr exists underlying ZrZO phase, which the authors  on  an  state  can  then exist as gaseous Zr(g) and ZrO(g). Nucleation and  growth of the oxycarbide phase were observed by TEM  to occur within carbon black agglomerates, which may  arise from condensation of ZrO(g) and CO(g) followed by  solid state diffusion to ZrCxOy(s). The oxycarbide is then  further reduced by contact with CO(g), carbon enriched particle until eventually the  leading to a more  carbide  is  formed. A schematic of the mechanism is shown in Fig. 1.  Recent  studies, however, have shown the dwell  time  and dwell temperature of the carbothermal reduction of ZrO2 to ZrC can be reduced using a solZgel type route. containing ZrZOZZr Typically, a polymer links is  formed from Zr n-propoxide or zirconium oxychloride,  and a  carbon source  such as  alcohols or  sugars  are  added. Pyrolysis then results in intimate mixing of ZrO2 and C particles, which have a higher reactivity and  therefore can be carbothermally reduced at peratures to coarse powders. Tao et al.13 reported the fabrication of ZrC at 1673 K with particle sizes of ,50-  lower  tem 100 nm using a sol derived from zirconium oxychloride. Xie et al.14 have also reported that a similar result from  starting materials of zirconium n-propoxide and sucrose  achieved 60-100 nm sized ZrC particles after reacting at  1973 K; however, with the added use of pulse current  heating,  the authors report a further reduction in reac tion time.  No studies  into the mechanism of  carbothermal  re duction and nitridation of ZrO2 are available in the literature as far as the authors are aware; however, the  nitridation  step  of ZrC has  been reported. Figure 2 al.15  shows  the  proposed  schematic  by Harrison  et  Reaction of ZrC powders  fabricated by carbothermal  reduction was  nitrided  under  10% doped H2-N2  at mosphere at 1800-1873 K for 4-24 h. During the initial  stages of nitridation, ZrC is destabilised to the  sub stoichiometric  carbide  (shown  in  step  1  of Fig.  2),  leading to the  formation of particles  containing edges  and terraces observed by SEM. The removal of carbon  leads to the formation of HCN(g) as observed previously by Bardelle and Warin16 in the carbothermic reduction-  nitridation of UO2. Nucleation of ZrCxNy (step 2 in Fig. 2) then occurs by  either evaporation-condensation of Zr(g) and HCN(g) at  the surface of the ZrC particle, although this is unlikely due to the high boiling point of Zr (,4500 K) resulting  in low amounts of Zr(g). More likely, nucleation occurs  at  the surface of  the ZrC1-x(s) particle from reaction of Zr with HCN(g) and proceeds by surface diffusion of Zr to the nucleation site. ZrCxNy particles then grow by  further migration of Zr  to the reaction layer and con densation of HCN(g). An external on the ZrCxNy particle as carbon diffuses to the reaction  layer is then formed  layer  to form more HCN(g),  leading to a carbon rich  external  layer and a nitrogen rich core (steps 3 and 4 in  Fig. 2). The reaction continues by further diffusion of  carbon  to  the  reaction  layer,  leading  to  removal  of  carbon at the surface until a core of ZrN exists. In the later reactions (.24 h, 1800 K), no monolithic  ZrC exists,  indicating that it has all been consumed and  the remaining particles are ZrN with an external shell of  ZrCxNy,  shown in step 5 of  the mechanism in Fig. 2,  1  Schematic of carbothermic reduction of ZrO2 reproduced from David et al.12  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  295  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'similar to the mechanism of nitridation of SiO2.17,18 The external carbon containing layer now exists as a solid  solution with the nitride.  Sintering  During  densification  of  non-oxide materials,  surface  oxides on the non-oxide particles have  contributed to coarsening.19  poor  densification  by  enhancing  grain  Spark plasma sintering (SPS) has been increasingly used  to prepare dense samples of nitrides. SPS, which is also  referred to as pulsed electric  current  sintering or field  assisted sintering,  involves  the passing of a pulsed DC  current  through the graphite die and the sample inside  heating via rapid Joule heating along with uniaxial pressing to compact powders.20,21 A review by Munir et al.22  reports  the success of SPS in achieving cleaner  grain  boundaries,  improved mechanical  and  thermo electric properties, densities close to theoretical density  (TD),  low grain growth, and improved oxidation and  corrosion resistances.  SPS offers the advantages of sintering materials in a  matter of minutes as opposed to the hours needed using  other conventional densification techniques such as hot  pressing or pressureless  sintering.  It  is  generally well  established that rapid Joule heating and current effects  in materials with good conductivity are mechanisms that al.23  operate during  the  sintering process. Hulbert  et  recently studied the mechanism of  ‘a momentarily gen erated’ spark plasma between particles. Plasma is a state  of matter similar to gas in which a certain proportion of  the particles are ionised. However, the authors noted the  absence  of  plasma  generation  and  suggest  that  the  enhanced densification of SPS lies in other factors such  as the rapid heating rates.  Sintering behaviour and microstructures of  the inert  matrix fuel (IMF) carbides and nitrides have been studied by Ryu et al.20 who used SPS to fabricate  IMF  pellets of actinide oxide phases in carbides and nitrides.  Dy2O3 was used to simulate minor actinide (MA) oxides. When comparing the sintering behaviour of the mono lithic  carbide  and nitrides with the Dy2O3 dispersion al.20  mixture, Ryu  et  found  that  the  sintering  onset  temperature was  lower  for  the simulated actinide con taining  composite. Density measurements were  per formed  using  the water  immersion method with  the  monolithic carbides and nitrides having densities of 63  and 65% TD respectively, and with addition of 20 wt-%  Dy2O3, 74% TD respectively. Perhaps other factors were causing  the TD for ZrC and ZrN achieved was 68 and  the high porosity of the materials, such as particle size or  sintering  programme,  as  these  values are less than doping.21,24,25 This  observed elsewhere without oxide  reduction in porosity was attributed to the removal of  surface oxide and enhanced diffusion as observed by others.26,27 Sakai and Iwata26 noted that  the self-diffu sion coefficient of AlN increases with oxygen content  due to cation vacancies introduced in AlN crystals by oxygen doping. Khor et al.27 reported that addition of  Sm2O3 mation of a liquid phase in Sm2O3-Al2O3-AlN system,  to AlN enhances AlN diffusion due to the for and this liquid phase increases mass transport. Hollmer28  reports  the manufacturing methods  for  (U, Zr)N fuels by SPS using both commercially available  and laboratory produced powders synthesised from pure  2  Schematic diagram of mechanism of nitridation of ZrC particles15  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  296  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'metal starting materials. Individual powders of UN and  ZrN were produced by a hydriding-nitriding of metal  ingots followed by mixing and SPS. Pellets of 99.7% TD  were  achieved  using  laboratory  fabricated  powders  compared with 89.7% TD using commercially available  ZrN. Heating rates and dwell  temperatures were kept  the same to compare densification (sintering parameters 21 and holding temperature of 1473 of 50 or 100 K min or 1923 K). Hollmer28 attributes this observation to the  high  oxygen  impurities  of  the  commercial  powder  (7.697 wt-% compared with 0.734 wt-% for the produced  powders),  although  no  experimental  technique  for  chemical analysis  is given and no errors are  reported.  Therefore,  it  is not appropriate  to discuss  the  results  from the commercial powder as a nitride because almost  equivalent amounts of N and O are present  (compo sition of 31.1 at-% N, 23.7 at-% O and 45.2 at-% Zr  assuming no presence of other elements). However, high  densities  (99.7% TD) were achieved using the as  fabri cated powders, suggesting that a small oxygen contamination (*0.7 wt-%) has  amount  of  little effect on  the achievable density.  Dysprosium nitride (used as a U/MA surrogate) and  zirconium nitride composites have been fabricated by Pukari and Takano29 by the hydrogenation-nitridation  of Dy and Zr metal ingots. Oxygen concentrations up to  1.2 wt-% were added using ZrO2 or by increasing milling  times to deliberately increase O content as to assess the  effect on the processing of these materials. Powders were  pressed into green bodies at 300 MPa and then pres sureless sintered at 1923 or 1973 K for 6 h. Pukari and Takano29  conclude  that oxygen content  that has dis solved into the nitride phase  increases  the achievable  densities  of  ZrN and  (Dy,Zr)N. Densities  of  ZrN  increased from 89.6 to 91.6 TD for 0.20 and 0.8 wt-% O  respectively, and (Dy,Zr)N increased from 93.1 to 94.5  TD for 0.26 and 0.59 wt-% O respectively. However, no  error ranges on the density measurements are given and  only geometrical densities are reported, and these details  would be necessary  to asses  the  significance of  these  changes.  The observations of increasing density with oxygen content by Pukari and Takano29 agree well with the high densities achieved by Hollmer.28 Low concentrations of oxides  that can be further oxidised on may act as oxygen getters,  reducing oxygen contamination of  the non-oxide particle  surface, reducing the mechanism of grain coarsening, and sintering.19 However,  allowing  increased solid state  the  formation of  secondary phase oxides will occur and the  effects on thermal properties may be detrimental.  Properties of ZrC, ZrN and ZrCN  Structure and bonding  Zirconium contains the electronic structure [Kr]5s24d2, and carbon and nitrogen have the structures [He]2s22p2 and [He]2s22p3  respectively. Carbon and nitrogen are  more  electronegative due  to their  increased ability to  attract  electron  density  arising  from poor  electron  shielding of  their nucleus. Nitrogen is one of  the most  electronegative atoms due to increased pull on its outer  electrons arising from an extra proton in the nucleus.  This difference of  electronegativity between the metal  and the non-metal  is what gives  rise to the ionic char acter of the bonds.  Both the carbide and nitride show a degree of covalent  bonding that arises  from interactions between the 2p  state  of  the  non-metal  and  4d  state  of  zirconium,  resulting  in metal-non-metal  as well  as metal-metal  bonding, but essentially with no non-metal to non-metal  bonding. Metallic  bonding  arises  when  atoms  are  ionised,  leaving the positive metal atoms and the delo calised electrons  that move  freely through the  lattice.  Electrostatic  forces  between  the  delocalised  electrons  and the positive metal atoms give rise to the bonds. The  high thermal and electrical conductivity of metals (and  the carbides and nitrides of some transition metals) arise  from the  ability of  the delocalised electrons  to move  through the lattice. The formation of these bands can be  described using molecular orbital theory shown, where a  diatomic molecule combines two atomic orbitals to form  two molecular orbitals.  In metals,  the electrons  in the  higher  energy  states  are  delocalised  over  the whole  crystal. As more molecular orbitals are formed, from the  enormous amount of atomic orbitals in a metal  lattice,  the energy difference between adjacent molecular orbitals decreases.30 The  gap between bonding  and anti bonding orbitals,  in the case of metals, also decreases,  leading to a continuum level. A density of electron states  arises  from the overlap of bonding  and antibonding  orbitals from the non-metal 2p orbital and the d metal  orbital., which has been reported to be just under level for the carbide.31  the  Fermi  The band structure of  the carbides and nitrides con tains a non-metal 2s, a non-metal 2p to zirconium d,s  and zirconium 4d,5s,5p states. There is a minimum in  the density of states (DOS) at the Fermi level when there  are 8 electrons  in the valence band, also known as a  valence electron concentration (VEC). At a VEC of 8,  the Fermi level is found between the non-metal p to metal d and d-like bands.1 With a VEC of w8, there are level,32 which is  a higher number of DOS at  the Fermi  located in the low energy region of et al.33  the metallic band.  Schwarz  calculated the DOS of ZrC and ZrN  using  augmented  plane wave  calculations, which  are  shown in Fig. 3.  Phase diagrams  Zirconium carbide and zirconium nitride are both crys talline face (Fm \\x163m,  centred  cubic  (FCC),  rock  salt  structure  space  group  225), whereby  the  carbon  and  nitrogen atoms occupy the  interstitial octahedral  sites  (shown in Fig. 4).  The carbide and nitride can accommodate non-metal  vacancies and are generally written with the non-stoi chiometric  formula ZrC1-x or ZrN1-x. The phase diagrams for ZrC an ZrN are shown in Figs. 5 and 6 respectively.34,35  Figure 5 reveals the existence of the ZrCx phase between *37.5 and 49.5 at-% C, where x is between 0.6  and 0.98, with the congruent melting of the ZrCx phase (roughly around ZrC0.85) at 3700 K. Below *37.5 at-% C, a hexagonal close packed (HCP) Zr and ZrCx two phase field exists until *1200 K, where  the Zr metal  becomes body centred cubic (BCC). This phase begins melting at *2127 K,  resulting  in liquid Zr and solid  ZrCx. It is also important to note that there is no existence of a ZrCx phase, where x is w1. Instead, ZrCx þ C phases form at carbon contents w49.5 at-%.  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  297  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  3  DOS of ZrC (left) and ZrN (right)  (Ef 5 Fermi  level) reproduced from Schwarz et al.33 showing higher DOS at Ef  for ZrN  The ZrZN phase  diagram (Fig.  6)  reveals  a  gas  phase, a liquid phase, the solution of nitrogen in BCCZZr (b-Zr) and the solution of nitrogen in HCP-Zr (a-Zr) as well as the non-stoichiometric FCC s-ZrNx phase, where x<1 at N at-% less that 50. Ma et al.36  reported a ZrN4 phase in the region of the single phase s-ZrNx; however, Lengauer32 stated the zirconium nitrogen system consisted of one nitride phase of sZrN1-x, similar to the carbide, where the N/Zr ratio <1. where x ¼ 0.5-1 below 50 at-% N similar to the ZrZC Gribaudo et al.35 also state the s-ZrNx phase exists system; however, above 50 at-% N, a solid ZrN phase  and N2 gas phase exist with no higher nitrides.  The ZrNx phase at  the Zr  rich boundary has been  examined  using  several  techniques,  such  as metallo graphy, X-ray diffraction (XRD), hardness and electrical resistivity. Eron’yan et al.37 used vapour pressures  and XRD to determine congruent melting of the ZrNx took place w3673 K phase at close to stoichiometry under a partial pressure of N2 of *6 MPa. Binder et al.38 also reported a tentative ZrZCZN  4  Rock  salt  crystal  structure: metal  atoms  (large green  spheres) and carbon/nitrogen atoms (small red spheres)  ternary diagram for  the  carbonitride at 1423 K using analysis of ZrZCZN  microprobe  and metallographic  samples prepared from mixing commercially available  ZrC and ZrN powders to achieve desired stoichiometries  (Fig. 7).  5  Calculated phase diagram for ZrZC system; reproduced  from  Fernandez  Guillermet34  from  Phase  Equilibria  6  Calculated phase diagram for ZrZN system; reproduced  Diagrams Online Database  (NIST Standard Reference  from Gribaudo et al.35  from Phase Equilibria Diagrams  Database  31),  the American Ceramic Society  and the  Online Database (NIST Standard Reference Database 31),  National  Institute of Standards  and Technology,  2015;  the American Ceramic Society and the National  Institute  figure number Zr-568  of Standards and Technology, 2015; figure number Zr-573  298  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'Figure 7 indicates that at 1423 K, the carbonitride has  a quasi-binary solid solution, which is due to the carbon  and nitrogen atoms occupying the octahedral interstitial sites randomly denoted in Fig. 7 as the d phase, whereby d-Zr(C,N)1-x(ss) denotes with this B1 rocksalt structure  the solid solution carbonitride (Fm \\x163m,  space  group  225). The metal  rich boundary is essentially a straight line dividing the d-Zr(C,N)1-x phase and the HCP a(ZrC,N) and BCC b(ZrC,N) phases. A three phase  equilibrium triangle  of  d þ a þ b  also  exists with  a  higher nitride composition of the Zr(CxN1-x)y phase.  Thermal and electrical properties  Enthalpy of  formation  Standard enthalpy of formation (DfH) for ZrC and ZrN at near stoichiometry is 2196.65 and 2365.26 kJ mol respectively.1,39 Toth40  21  summarised  the  heat  of  for mations  for  ZrCx  phases  with  increasing  carbon  vacancies  having  to  be  less  negative, which  can  be  explained as the ZrC bond strength decreasing, resulting  in less  stable compositions. This difference in enthalpy  of formation between ZrC and ZrN matches well with Pierson1 comments that the ZrN bond is stronger due to  the increased electronegativity difference nitrogen and carbon. However, Lengauer32  between  states  the  nitride bond is weaker than the carbide due to the extra  electron in the nitrogen 2p orbital occupying an antibonding orbital. Lengauer32  also explained the more  negative the heat of formation for the nitrides, the lower  the equilibrium nitrogen pressure is, which determines  the phases obtained, and if the pressure is high enough, ratio w1. However, this leads to a non-metal/metal ZrNx metastable phases with xw1 (such as Zr3N4) are usually only produced using magnetron sputtering or  chemical vapour deposition techniques to produce thin  films, where possible.4-6,41  observation  of  metastable  phases  is  Heat capacity  Heat  capacity  (Cp) energy added to a material to the resulting temperature  corresponds  to  the  ratio  of  heat  change and shows how much energy is required to heat a  mole of material by 1 K. Heat capacities for ZrC plotted as a function of temperature are shown in Fig. 8.42-46 Heat capacities of ZrC were measured by Neel et al.45 and Levinson44  using  high  temperature  drop  calori metry.  Scatter  in  the  literature  is  due  to  authors  reporting different stoichiometries of ZrCx, where x can be between 0.92 and 0.96. Chase39 reported the function  of Cp against temperature (where Cp is heat capacity in J mol and T is the absolute temperature) by  21 K  21  equation (3).  C pZrC ¼ 51:027 þ 3:685T 2 0:199T 2 þ 0:028T 3 2 1:304=T 2  ð3Þ  Molar heat capacities for ZrN and some zirconium carbonitride phases are shown in Figs. 9 and 10.21,45,47-49  Differential  scanning  calorimetry  (DSC) was  used to measure heat capacity of ZrN by Basini et al.,47 Ciriello et al.,49 and Muta et al.21 and by drop calorimetry  7  Phase diagram for ZrZCZN system at 1423 K reproduced  from Binder  et  al.38  from Phase  Equilibria Diagrams  Online Database (NIST Standard Reference Database 31),  the American Ceramic Society and the National  Institute  of Standards and Technology, 2015; figure number Zr-577  8  Heat capacities of ZrC as function of  temperature (pro tective atmosphere conditions are given in legend where  available)39,42-46  9  Heat capacity of ZrN as function of  temperature (protec tive atmosphere conditions are given in legend where  available)21,39,45,47-49  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  299  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'by Hedge et al.48 and Neel et al.45 Chase39 reported the  function of Cp against temperature in equation (4).  C pZrN ¼ 46:4587 þ 7:0063T 2 0:0077T 2 þ 0:0007T 3 2 0:7205=T 2  ð4Þ  Both the carbide and nitride show a sharp increase in Cp at low temperatures to ,600 K, then plateauing until 1500 K, where they begin to increase again. Nitrides  generally show a slightly higher heat capacity than car bides;  however, with experimental ,42 J mol  room temperature  (RT)  values  of  21 K  21  compared  to  38 J mol  21 K 21  for  ZrN and  ZrC respectively,  the  differences can fall within experimental error and so can  be considered negligible. Lengauer et al.50 measured the heat capacities of some  carbonitrides of zirconium using DSC up to 1100 K; again, Cp increased with increasing nitrogen content.50 Einstein’s formula for predicting molar heat capacity is  derived from the oscillation of the bonds of a molecule,  explaining that as T approaches 0 K, heat capacity also  reaches zero and the plateau of Cp observed in materials  at higher  temperatures arises  from saturation of oscil lating bonds with energy. From this  explanation,  the  higher Cps of  the higher nitrogen containing carboni trides arise from the nitride containing more occupied  antibonding orbitals, giving longer, weaker bonds  that  are less stiff. This results in the vibrational energy levels  being  closer  together;  therefore, when  a  quantity  of  energy arrives as heat,  there is  little change in the dis tribution  in  the  population  of  the  energy  levels  (according  to the Boltzmann distribution)  that  corre sponds to a material with a higher heat capacity.  Vapour pressure  The vapour pressure of ZrC has been more extensively  studied than ZrN, with the former having a reported vapour pressures of Zr(g) at 2600 K of 6|10 24 Pa and 7|10 24 for Knudsen and Langmuir measurements respectively.24,51,52 The vapour pressure of Zr(g) over ZrN at 2500 K is *3|10 23 Pa.53 The higher vapour press ure  of  the  nitride compared with reported by Pollock51 can be explained due stronger nature of the ZrZC covalent bond.  the measurement  to  the  Thermal conductivity  Thermal properties, such as thermal conductivity, are an  important property for nuclear fuel elements. The desire  for  increased  efficiency  for  generation  IV  reactors  necessitates high thermal  conductivity  to increase  the  rate of heat transfer from the fissile phase to the coolant.  Safety is another paramount  concern for all materials  being evaluated for use  in generation IV NPP.  IMFs  must  reduce the potential  for  the buildup of hot  spots  within the reactor core, where fission events and their  daughter decay events will cause concentrated hot spots  in the event of a loss of coolant accident event.  Thermal conductivity (W m the thermal diffusivity (a, m2 s  21 K 21)  is the product of 21), density (r, kg m 21) of a material, given  23)  and heat capacity (Cp, J kg by equation (5).  21 K  K ¼ arC p  ð5Þ  Thermal conductivities can be corrected to represent  the thermal conductivity of a material having 100% TD  using the Maxwell-Eucken equation (equation 6), where  P is  the porosity, Kp is  the measured thermal conduc tivity and KTD is the corrected thermal conductivity for a fully dense material.54  K p ¼ ð1 2 PÞ  ð1 þ PÞ K TD  ð6Þ  Thermal conductivity of ZrC and ZrN as a function of  temperature has been studied extensively (Figs. 11 and 12).21,25,32,45,47-49,55-60 All values in Figs. 11 and 12 have  been corrected to 100% TD for  ease of  comparison;  however, et al.47  the  experimental  values obtained by Basini  for a 70% TD ZrN sample are shown as  they  match quite well with literature, and extrapolation to a  fully dense material using equation (6) may be inappropriate for such a large porosity.47 A positive trend is  seen for both the carbide and nitride, which is evidence  for metallic bonding. Thermal conduction in ceramics is  a contribution of phonon and electron conduction.  In  defect free materials, phonon conduction increases with  increasing temperature as phonon-phonon collisions are  rare,  until  a  point where  lattice  vibrations  become  anharmonic  causing  scattering  of  the  phonons  and  11  Thermal conductivity of ZrC as function of  temperature  (protective atmosphere conditions are given in legend  where available)32,55-57,60  10  Heat  capacity  of  some  zirconium carbonitrides  as  function  of  temperature measured  under  Ar  from  Lengauer et al.50  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  300  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'reducing thermal conductivity.  It  is  therefore expected  that  if  the thermal conductivities of ZrC and ZrN were  wholly dominated by phonon conduction,  there would  be a decrease at higher temperatures; however, this is not  observed in Figs. 11 and 12, with the exception of some work on ZrN reported by Basini et al.,47 Hedge et al.,48 and Neel et al.,45 where thermal conductivity decreases  at higher temperatures.  The carbide and nitride of zirconium can accommo date vacancies, particularly on the non-metal  site, and  with regards to thermal properties, vacancies and point  defects  cause phonons  and electrons  to be  scattered,  which decrease thermal conductivity. The effect of non stoichiometry on the  thermal  conductivity of ZrC has  received attention, with the thermal conductivity falling to ,10 W m  rapidly  21 K  21  for  ZrC0.5-ZrC0.85.24,61  However, no study of the effect of thermal conductivity  as a function of nitrogen vacancies exists for the nitride.  This may be the reason for the large amount of scatter in  the literature as vacancies,  impurities and porosity will  all adversely affect  thermal conductivity, with the stoi chiometry rarely being reported in the literature. Comal.60 were  mercial ZrN powders used by Harrison et  reported to have a stoichiometry of ZrN0.86 and agreed with much of the published Indeed, the more recent studies of the thermal conductivity of ZrN59 and high N containing ZrCN ceramics60 have  data.45,47-49  shown  thermal conductivity values of a stoichiometric ZrN to be ,55 W m 21 at RT (Figs. 12 and 13  phase  21 K  respectively).  Lengauer  et  al.50  and Harrison et  al.  reported the  thermal  conductivities of  several ZrCxNy  ceramics  to  1100 and 2073 K respectively, al.50  shown in Fig. 13. Len gauer  et  produced  the  desired  compositions  by  mixing of  commercially available ZrC and ZrN pow ders, and chemical analysis was performed by combustion gas analysis. Harrison et al.60 prepared the ZrCxNy powders via carbothermic reduction-nitridation, and  stoichiometries are reported from combustion gas anal ysis. The trend in increasing thermal conductivity with  increasing nitrogen content was attributed to a higher  number of  electrons  contributing  to thermal  conduc tivity with increasing nitrogen content as determined by  electrical  conductivity measurements  (presented in the  following section).  Electrical conductivity  Thermal  conductivity  is  a  sum of  the  electron  and  phonon conduction of a material expressed in equation  (7), where KTot ductivity, Kel (W m and Klat (W m  (W m  21 K 21) 21 K 21) 21 K 21)  is  the total  thermal con is the electron contribution  is  the phonon contribution to  thermal conductivity. The electronic contribution can be  determined  using  the  Wiedemann-Franz-Lorenz  equation (equation (8)), where L is the Lorenz number, widely accepted as 2.44|10 28 WV K 22, s is (V21 m  the elec trical conductivity temperature (K).25  21)  and T is  the  absolute  K Tot ¼ K lat þ K el  ð7Þ  K el ¼ LsT  ð8Þ  The Wiedemann-Franz law states that for metals, the  ratio of thermal conductivity to electrical conductivity is constant.62  proportional Morgan63  to  temperature  and  is  a  reported that  the  electrical  contribution to  thermal conductivity of good electrically conductive (>103 V21m 21<105 V21 m 21) ceramics obey this law of  proportionality, and their electrical heat conduction can  be  determined  using  the Wiedemann-Franz-Lorenz  equation (equation (9)). Owing to the phonon conduc tion still being an important factor in these materials, the  Wiedemann-Franz-Lorenz  equation corresponds only  to the Kel contribution of KTot.  In the case of metals,  their  thermal conductivity is completely dominated by  electrical heat conduction, Lorenz equation ,KTot.  so the Wiedemann-Franz-  K  s  ¼ LT  ð9Þ  Electrical conductivities of ZrC and ZrN as a function temperature are presented in Fig. 14.25,56,58,64-67 The  of  electrical conductivities of some zirconium carbonitride  phases are shown in Fig. 15, although no values for the monolithic phases were presented in the same work.50  ZrC and ZrN both show metallic behaviour, with an  inverse  relationship  to  temperature. Factors  such  as  porosity,  impurities, vacancy concentration and micro structure will  affect  the  electrical  conductivity of  the  samples, which is again the cause of  scatter  in the lit erature, mainly  for  the ZrN samples. However,  no  12  Thermal conductivity of ZrN as function of  temperature  (protective atmosphere conditions are given in legend  where available)21,25,32,45,47-49,58-60  13  Thermal conductivities of ZrCxNy phases as function of  temperature from Lengauer et al.50  (under high purity  Ar) and Harrison et al.60 (under high purity Ar)  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  301  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'detailed work has been reported on how defects, oxygen  and  carbon  impurities  and  microstructure  affect  conductivity. Adachi et al.25 measured the electrical conductivity of ZrN samples containing ,10-18% porosity, determining  nitrogen content using chemical analysis to give a stoichi ometry of ZrN0.935; however, no oxygen or carbon analysis is presented. Vacancies and impurities are an important  factor in conduction, as defects create a scatter point for  phonons and electrons, reducing the electrical conductivity. Adachi et al.25 also reported the electron contri bution to thermal conductivity (equation (8)) as 16 and 21 for the 18% and 10% porous samples re 20 W m  21 K  spectively, which amounts to two-thirds of the total conreported. Petrova et al.67 reported the RT conductivity of ZrN as ,4506104 V21 m  ductivity  electrical  21  using samples produced by precipitation of ZrCl4, with N2/ H2 gas achieving stoichiometries of between ZrN0.97-1.00 and sample densities very close to TD, stating that this  preparation route seems to lead to high purity materials;  however, no experimental technique for elemental analyses  is presented.  It can be seen from Figs. 14 and 15 that the nitride has  a higher electrical conductivity than the carbide, and for  the carbonitride, electrical conductivity increases with increasing nitrogen content. Lengauer32 attributed this  increase in conductivity to the nitride having a higher 3),33  DOS  at  the Fermi Level  (Fig.  and  due  to  the  nitrogen’s  extra  valence  electron resulting  in an anti bonding state,  the  symmetry of  the p-d non-metal  to  metal  bond  changes,  allowing  increased metal  d-d  bonding,  resulting in this  increased DOS at  the Fermi  Level.  Thermal expansion  Interatomic spaces between metal and non-metal atoms  can vary as a function of  temperature due to increased  energy exciting the atoms into higher energy states. This  causes higher  amounts of  vibration, which results  in  them moving  further  apart  from each  other  due  to  anharmonic  behaviour  when  oscillating.  The  bond  strength affects the extent to which the atoms can move  away  from each other, with stronger bonds  reducing  thermal  expansion  and  weaker  bonds  showing  the  opposite  effect. Thermal  expansion  is  an  important  character  of  a  composite  fuel material,  such  as UN  dispersed  in  ZrN.  If  the  two materials  expand  by  different amounts on heating,  there is a thermal expan sion mismatch, which can result  in the ceramics crack ing. This is a potential  issue for advanced nuclear fuels,  as cracking can lead to the release of fission products  from the pellet. The thermal expansions of ZrC and ZrN neutron diffraction,68 dilatometry25,45,58,71  measured  by  high  temperature  XRD69,70  and  are  presented  in  Figs. 16 and 17. Adachi et al.58 explained that thermal expansion is not  affected by porosity and grain size, accounting for  the  similar  results  for  their porous  samples with previous  work. For  these materials,  the  expansion in length is  expressed as the coefficient of thermal expansion (CTE), expansion, DL, as a unit K 21. The  which  represents  the  degree  of  function  of  temperature  and  has  the  carbides and nitrides presented in Figs. 16 and 17 all range of 2-8|10  have CTEs  in the  26 K 21  from RT 1273 K and are similar to other ceramics such as alumina and zirconia. Aigner et al.69 measured the var iance in thermal expansion of ZrC, ZrN and ZrC0.5N0.5  using high temperature XRD, giving the CTEs up to 1600 K (aav)  in equation (10) and giving CTEs for 7.5,  16  Elongation with respect  to length at 298 K of ZrC as  function  of  temperature  (protective  atmosphere  con ditions are given in legend where available)45,68,71  14  Electrical conductivity of ZrC and ZrN as function of  temperature  (protective  atmosphere  conditions  are  given in legend where available)25,55,56,58,64-67  15  Electrical  conductivities  of  ZrCxNy  phases  at RT  as  function of  [C]/[C] 1 [N] ratio from Lengauer et al.50  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  302  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', '7.65 and 7.8|10  26 K 21  for ZrC, ZrC0.5N0.5 and ZrN  respectively.  aav ¼  7:8 þ 0:3  ½C \\x8a ½C \\x8a þ ½N \\x8a ½C \\x8a ½C \\x8a þ ½N \\x8a  \\x12  \\x13  \\x12  \\x13  2 0:6  \\x12  \\x13  £ 10  26  ð10Þ  Mechanical properties  ZrC and ZrN have received extensive use as coatings for  cutting tools and abrasives due to their high hardness  and melting points. As a consequence,  the majority of  work regarding the mechanical properties of ZrN and  ZrC, such as Young’s modulus and hardness, has been performed on thin films and coatings.72-81  Young’s modulus  Young’s modulus of ZrC as a function of temperature was measured by dynamic methods by Baranov et al.79 and Travushkin et al.80 and by four-point bend method al.82  by  Zubarev  et  (Fig.  18). Measuring  strain  in  materials with  high Young’s modulus  values  (in  the  regions GPa) with high tensile  strengths  (in the MPa  region)  via  static methods  is difficult due  to the  low  values strains; thus, measurements from vibrational data  Fig. 18 are likely more accurate than the bending data. Desmaison-Brut83  Alexandre  and  measured  the  Young’s modulus of ZrN, which was produced using hot  isostatic pressing at 2223 K under 195 MPa of pressure  for between 1-2 h. Elasticity measurements were made  by the disc vibration technique and reported as a func tion of porosity, with highly dense (1% porosity) samples having a Young’s modulus of *390 GPa, decreasing to  290 GPa with increasing porosity (8% porosity). Adachi et al.84 measured the Young’s modulus of ZrN ceramics  prepared by SPS using ultrasonic pulse  echo method.  Samples of  varying porosity  (7-18%) were measured,  giving values of 288 GPa for the 8% porous sample and 156 GPa for the 18% sample. Yang et al.85 measured the  Young’s modulus  of  zirconium carbonitrides  by  the  ultrasonic  pulse  technique,  reporting  a  variance  of  390 GPa for a composition of around ZrC0.1N0.9 and reaching a maximum of 420 GPa for a composition of  ZrC0.8N0.2, which decreases to 390 GPa for monolithic  ZrC. The  increase with small amounts of nitrogen is  attributed to the sigma bonding states between the non metal p orbital and metal d orbital being filled with the  increased valency arising from nitrogen bonds with a  maximum bond at a VEC of 8.4. However,  this  then  decreases with increasing further nitrogen content.  Hardness  ZrC has  received many more  hardness  studies  than  nitride (Fig. 19), with indentation being used to study bulk hardness and preferred slip systems.86 The covalent  bonding in ZrC means that in closer packed planes, the  bonds will be stronger than in a metal, giving a harder  material.  Hardness as a function of died by Kohlstedt87 on single crystals of *1273 K. Discs of 10 mm in diameter and 0.5 mm in  temperature has been stu ZrC0.94  height were orientated in the (001) plane and placed under a load of 500 g at a rate of 10 mm s 21. Kohlstedt87  noted that at lower temperatures,  the highly directional  covalent bonds on the close packed {111} planes inhibit  slip and reduce dislocation mobility. However, at higher  17  Elongation with respect  to length at 298 K of ZrN as  function  of  temperature  (protective  atmosphere  con ditions are given in legend where available)25,58,69,70  18  Young’s modulus of ZrC as  function of  temperature  (protective atmosphere conditions are given in legend  where available)79,80,82  19  Hardness of ZrC as function of  temperature (protective  atmosphere  conditions  are  given  in  legend  where  available)87,89,90  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  3 0 3  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'temperatures,  sharp  decreases  in  hardness  were  explained due to the increasing mobility of s electrons,  with temperature  screening  the bonding orbitals  that  reduce the directionality of  the covalent bonds,  thereby  reducing the hardness to that similar of a metal. How ever, this may be described more simply by the fact that  dislocations can only move if  they overcome resistance  of the lattice, which will have to overcome an activation  barrier, which  is  normally  not  high; thus, energy, kT, can have a significant impact.88 al.89 measured the microhardness of  thermal  Kumashiro et  single  crystals of ZrC0.9, NbC0.9 and TaC0.83 using a load of 100 g for varying loading times (1-100 s), with  the indentation orientation in the {100} plane in the v001w direction. ZrC0.92 was found to be the softest, justified by ZrC having a larger lattice parameter than  the other  carbides,  reflecting  the metal  carbide bond  being longer and therefore weaker. Thermal  softening  was observed as it was by Kohlstedt, which was attrib uted to increasing lattice vibrations,  resulting in bond  lengthening and weakening. Gridneva et al.90 measured the Vickers hardness (HV) of  single crystal ZrC on the (100) plane under a load of 115 g  to a temperature of 1200 K. The authors noted a sharp  reduction in the number of microcracks around the point of  indentation at 1123 K, at which temperature ZrC can be  deformed plastically. Cracks produced from indentations spread in the v100w direction, with the {111} v110w and {110} v110w slip systems being active.  Microhardness of ZrN produced by SPS was measured by Adachi et al.84 using an HV tester at RT, with loads  between 0.98 and 9.8 N being applied for 15 s. HV was  characterised as a function of porosity, with denser samples  of 8% porosity having an HV of between 10 and 11 GPa  (depending on load), which decreased to between 5 and  8 GPa for more porous samples (18%). Results were ex trapolated to a fully dense sample, estimating an HV of *15 GPa for the samples. al.50 measured  Lengauer  et  the RT HV of  some  ZrCxNy phases using a load of 0.98N. HV was found to  decrease with nitrogen content, from a value of 23 GPa for ZrC to *18 GPa for ZrC0.25N0.75. The nitrogen on the hardness of the materials were believed  effects of  due to increased antibonding states in the nitride arising  from its extra valance electron. Although the nitride will  have more antibonding states occupied than the carbide,  decreasing the covalent nature of its bonding, the nitride  bond is also electrostatically stronger due  to nitrogen  being more electronegative than carbon.  Oxidation behaviour  Diffusion of atoms  in a material  is  theoretically gov erned by the existence of an atomic concentration gra dient whereby  atoms  diffuse  from an  area  of  high  concentration  to  low concentration.  Ideal  diffusion,  such as  that of a gas  through a porous  solid,  can be  described by Ficks first  law (equation (11)), where F is  the flux of atoms, D is the diffusion coefficient and dn/dx  represents the change in gradient in concentration in the  x direction.  F ¼ 2D  dn  dx  ð11Þ  However, another constraint that becomes important  in the diffusion of atoms in a solid is the energy needed  for that atom to move, known as the activation energy,  Ea. According to the Boltzmann distribution, tribution of atoms that populate higher energy states  the dis decreases exponentially and exp(2DE/kT), where DE (J)  so  is  proportional  to  is  the  energy  difference  between two states, k is the Boltzmann constant (1.381610 21) and T is the temperature (K). The diffusion coefficient, D (m2 s  223 J K  21)  is  related to this acti vation energy by equation (12), where D0 is a constant plot of lnD versus 1/T (T ¼ ‘). for that system and corresponds to the y intercept on a D ¼ D0 expð2E a =kT Þ  ð12Þ  The standard enthalpy of formation (DfH 0) of ZrO2 is 21097.46 kJ mol 2196.65 as compared to and 2365.26 kJ mol for ZrC and ZrN respectively.1,39  21  21  With little change between the entropy (S) of the oxide,  carbide and nitride (between 33 and 50 J mol  21 K  21),39  the reaction is driven by this large decrease in enthalpy.  Therefore, there is a large decrease in the free energy of  the system according to equation (13), where G is  the  Gibbs  free  energy, H is  the  enthalpy, T is  absolute  temperature and S is the entropy.  DG ¼ DH 2 T DS  ð13Þ  Zirconium nitride  The majority of oxidation studies of ZrN have been per formed on thin films due to the extensive use of ZrN as a hard coatings of tools.91-93 Krusin-Elbaum and Wittmer92  studied the oxidation kinetics of ZrN thin films, reporting the activation energy of oxidation to be 241+ 10 kJ mol  21  in the temperature range of 748-923 K with the oxide layer  comprising of monoclinic and cubic ZrO2. The authors  oxidised thin films with a thickness of between 60 nm and 1 mm under flowing O2. Krusin-Elbaum and Wittmer92 found non-linear kinetics with parabolic rate behaviours  being observed and state  that  the diffusion of oxygen  through the oxide layer  is  the rate limiting step. Cubic  zirconia was determined to be present as well as monoclinic  using glancing angle XRD; however, no comment is made  of the existence of this cubic polymorph or the effect it may  have on oxidation rate. Panjan et al.93 studied the oxidation of several metal 3 mm)  nitride  thin  films  (between  300 nm and  under  flowing O2 showed the oxidation of ZrN coatings to have an activation energy of 229 kJ mol 21 and also obey a parabolic  between  773  and  1123 K.  The  authors  rate law, agreeing well with Krusin-Elbaum and Wittmer.92 Both studies show a plateau region after oxidation, and Krusin-Elbaum and Wittmer92  initial  suggest  that  the diffusion of oxygen through the oxide layer is  the rate limiting step. Electrical resistivity measurements  were performed on CrN thin films, which showed a  rapid increase in resistivity in which the authors attrib uted to oxidation of grain boundaries during the initial  stage of oxidation. Typical grain sizes of the nitride thin  films were 20-40 nm, and so it would be plausible that  oxidation of grain boundaries would result  in increased  electron scattering; however, no comment or compari son of the electrical resistivity increase of thin films with  temperature is given. et al.94 studied the oxidation of 30 mm ZrN  Caillet  coatings  in  the  range  of  823-973 K observing  linear  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  3 0 4  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c', 'kinetics, contrary to Krusin-Elbaum and Wittmer92 and Panjan et al.,93 suggesting that oxygen diffusion through  the oxide  layer does not  affect  the  rate. However, a at *823 K. The  significant mass  loss was  observed  authors propose HCPZZr at  that nitrogen is  initially lost  forming  the surface before oxidation of ZrO2. The  oxide  layer was  composed  of monoclinic  and  cubic  ZrO2, begins by destabilisation of ZrN, resulting in a Zr rich then becomes a saturated a-Zr  and  the  authors  propose  that  the mechanism  phase that  layer before  forming an oxide scale. Harrison and Lee95  studied the oxidation behaviour  of ZrN ceramics under  static  air over  a  temperature  regime  of  973-1373 K.  The  authors  proposed the low temperature (<1073 K) and (>1173 K) oxidation of ZrN. The temperatures <1073 K can be  two  mechanisms  for  higher  temperature  oxidation mechanism at  described  in  several  steps;  initially,  oxygen  diffusion  through the surface ZrN grain boundaries  results  in a  ZrOxNy phase, which then forms m-ZrO2. Stresses from  the m-ZrO2 intergranular fracture and the creation of a porous oxide  grains between the ZrN grains  result  in  layer, similar to the observation by Berkowitz-Mattuck.96 Oxidation of ZrN at higher temperatures, how ever, proceeds by the  formation of dense oxide  layer,  which was identified to be submicrometre c-ZrO2 grains  revealed by TEM and selected area diffraction patterns.  This high temperature polymorph is  stabilised by the  substitution of oxygen with nitrogen forming ordered  oxygen vacancies. This dense oxide layer  inhibits oxy gen’s ingress and therefore rate of oxidation. However,  at higher temperatures, extensive oxidation of  the ZrN  grains  is observed,  indicating that  this protective layer  fails as diffusion rates increase.  Zirconium carbide  ZrC has similar chemical properties and crystal structure to  ZrN, and its oxidation kinetics and mechanism in the bulk are better understood24,86,97-99 with onset of oxidation occurring between 653 and 750 K86 similar to that of ZrN.92 The mechanism of oxidation of ZrC at temperatures w743 K is explained with the following steps:97  (i)  Formation of  a ZrOxCy phase  surface, which  proceeds to form amorphous ZrO2 and C. Crystallites of cubic zirconia (c-ZrO2) nucleate,  (ii)  which form a dense oxide layer with free carbon  stabilising the c-ZrO2 polymorph (via substitution of O2- with C3- ions, creating oxygen vacancies).  (iii)  Oxygen then diffuses through this oxide layer to the  free carbon producing CO2, which diffuses out via any existing cracks or pores, leaving behind voids  and pores in the zirconia layer, which with very little  carbon left  to stabilise transforms  to monoclinic  zirconia (m-ZrO2), along with small amounts of tetragonal polymorph (t-ZrO2).86 The oxygen diffusion through ZrO2 has been reported as *10 21 at 1723 K,100 which is *10 times faster temperatures,101  210 m2 s  than for SiO2 effectiveness of limited.102  at  similar  and so the  the dense ZrO2  layer  is expected to be  Conclusions  This  article  reviews  the  current  available data for  the  properties of ZrC, ZrN and mixed carbonitrides phases  and identities causes of scatter in the literature and areas  requiring  further  research. The  processing  routes  and  thermal, mechanical and irradiation properties of ZrC,  ZrCN and ZrN materials have been surveyed. ZrC and  ZrN are both promising candidates for use in high tem perature environments, namely,  in next generation NPP,  due to their high melting temperatures, superior thermal  conductivity, good mechanical properties and increasing  positive  irradiation  results. However,  this  review has  highlighted some areas where data are largely scattered or  limited, with impurities, vacancies and microstructure all  affecting reported values. Fabrication experience of pow ders is limited, with the majority of thermal, mechanical  and irradiation properties being carried out on commer cially available ZrC and ZrN powders with quantitative  analysis of ceramics rarely reported. It  is clear that pro duction of pure, stoichiometric, dense ceramics is not yet  fully achieved, hindering the reporting of accurate ther mophysical data. 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Ettmayer:  ‘Lattice  parameters  and thermal  expansion of Ti(CxN1-x), Zr(CxN1-x),  Hf(CxN1-x) and TiN1-x from 298 to 1473 K as investigated by high-temperature X-ray diffraction’, J. Alloys Compd, 1994, 215,  (1-2), 121-126.  70. M. Takano, M. Akabori, Y. Arai  and K. Minato:  ‘Thermal  expansion of TRU nitride  solid solutions as  fuel materials for 389,  transmutation of minor  actinides’,  J. Nucl. Mater,  2009,  (1), 89-92.  71. B. A. Fridlender and V. S. Neshpor:  ‘Thermal expansion of  the  carbides of group IV-VI Metals (Transition Class)’, High Temp, 1976, 14,  (5), 847.  72. Y-C. Chieh, W-Z. Lo and F-H. Lu: ‘Microstructure evolution of ZrN films annealed in vacuum’, Surf. Coat. Technol., 2006, 200,  (10), 3336-3340.  73. C.  J. Tavares, L. Rebouta, M. Andritschky  and  S. Ramos:  ‘Mechanical characterisation of TiN/ZrN multi-layered coatings’, J. Mater. Process. Technol, 1999, 92-93, 177-183.  74. V. Craciun, E. J. McCumiskey, M. Hanna and C. R. Taylor: ‘Very  hard ZrC thin films grown by pulsed laser deposition’, J. Eur. Ceram. Soc, 2013, 33,  (1-2), 2223-2226.  75. H-M. Tung,  J-H. Huang, D-G. Tsai, C-F. Ai  and G-P. Yu:  ‘Hardness and residual stress in nanocrystalline ZrN films: effect  of bias voltage and heat A500, (1-2), 104-108.  treatment’, Mater. Sci. Eng. A, 2009,  76. D. Craciun, G. Socol, N. Stefan, G. Dorcioman, M. Hanna,  C. R.  Taylor,  E.  Lambers  and V. Craciun:  ‘The  effect  of  deposition atmosphere on the chemical composition of TiN and  ZrN thin films grown by pulsed laser deposition’, Appl. Surf. Sci, 2014, 302, 124-128.  77. D. Ferro, J. V. Rau, V. RossiAlbertini, A. Generosi, R. Teghil  and S. M. Barinov:  ‘Pulsed laser deposited hard TiC, ZrC, HfC  and TaC films on titanium: hardness and an energy-dispersive 202,  X-ray  diffraction  study’, Surf. Coat. Technol,  2008,  (8),  1455-1461.  78. Y. Long, A. Javed, J. Chen, Z.  -K. Chen and X. Xiong:  ‘Phase  composition, microstructure and mechanical properties of ZrC  coatings produced by 2014, 40,  chemical  vapor deposition’, Ceram.  Int,  (1), 707-713.  79. V. M.  Baranov, V.  I. Knyazev  and O.  S. Korostin:  ‘The  temperature dependence of the elastic constants of nonstoichiometric zirconium carbides’, Strength Mater, 1973, 5,  (9), 1074-1077.  80. G.  G.  Travushkin,  V.  I.  Knyazev,  V.  S.  Belov  and  G. A. Rymashevskii: ‘Temperature threshold of brittle failure in interstitial phases’, Strength Mater., 1973, 5, (5), 639-641.  81. E. Torokk, A. J. Perry, L. Chollet and W. D. Sproul:  ‘Young’s  modulus of TiN, TiC, ZrN and HfN’, Thin Solid Films, 1987, 153,  (1-3), 37-43.  82. P. V. Zubarev, A. B. Kuraev, A. S. Maskaev, E. V. Astakhova  and A. A. Bessarabov:  ‘Stress  relaxation in zirconium carbide.  Report 1. Test method and stress Mater., 1994, 26,  relaxation curves’, Strength  (1), 63-69.  83. N. Alexandre and M. Desmaison-Brut:  ‘Mechanical properties of  hot isostatically pressed zirconium nitride materials’, J. Mater. Sci., 1993, 28, 2385-2390.  84.  J. Adachi, K. Kurosaki, M. Uno and S. Yamanaka:  ‘Porosity  influence  on  the  mechanical  properties  of  polycrystalline 358,  zirconium nitride  ceramics’,  J. Nucl. Mater,  2006,  (2-3),  106-110.  85. Q. Yang, W. Lengauer, T. Koch, M.  Scheerer  and  I.  Smid:  ‘Hardness and elastic properties of Ti(CxN1-x), Zr(CxN1-x), Hf(CxN1-x)’, J. Alloys Compd, 2000, 309, (1-2), L5-L9.  86. Y. Katoh, G. Vasudevamurthy, T. Nozawa  and L. L. Snead:  ‘Properties of zirconium carbide for nuclear J. Nucl. Mater, 2013, 441, (1-3), 718-742.  fuel  applications’,  87. D. L. Kohlstedt:  ‘The temperature dependence of microhardness 8,  of  the  transition-metal  carbides’,  J. Mater. Sci,  1973,  (6),  777-786.  88.  J. Wang and L. J. Vandeperre:  ‘Ultra-high temperature ceramics:  materials  for  extreme  environment  applications’  ‘Deformation  and hardness of UHTCs as a function of  temperature’, 236-266;  2014, Hoboken, NJ, John Wiley & Sons Inc.  89. Y. Kumashiro, Y. Nagai  and H. Kato: ‘The Vickers microcrystals up to 15008C’,  hardness of NbC, ZrC and TaC single J. Mater. Sci. Lett., 1982, 1, (2), 49-52.  90.  I.  V.  Gridneva,  Yu.  V. Mil’man,  G.  A.  Rymashevski,  V.  I. Trefilov and S.  I. Chugunova:  ‘Effect of  temperature on  the strength characteristics of zirconium Metall. Met. Ceram, 1976, 15, (8), 638-645.  carbide’,  Powder  91. Z. T. Wu, Z. B. Qi, W. F. Jiang, Z. C. Wang and B. Liu: ‘Influence  of niobium addition on microstructure, mechanical properties and  oxidation resistance of ZrN coatings’, Thin Solid Films, 2014.  92. L. Krusin-Elbaum and M. Wittmer: ‘Oxidation kinetics of ZrN thin films’, Thin Solid Films, 1983, 107, (1), 111-116.  93. P. Panjan, B. Navinsek, A. Cvelbar, A. Zalar and I. Milosev:  ‘Oxidation  of TiN, ZrN, TiZrN, CrN, TiCrN and TiN/CrN  multilayer hard coatings reactively sputtered temperature’, Thin Solid Films, 1996, 281-282, 298-301.  at  low  94. M. Caillet, H. F. Ayedi and J. Besson:  ‘Etude de la corrosion de  revetements  refractaires  sur  le  zirconium i.  Oxydation  par  l’oxygene de revetements de nitrure Common Met, 1977, 51, (2), 305-322.  de  zirconium’,  J.  Less 95. R. W. Harrison and W. E. Lee:  ‘Mechanism and kinetics of  oxidation of ZrN ceramics’, J. Am. Ceram. Soc,  in press 2015.  96.  J.  B.  Berkowitz-Mattuck:  ‘High  temperature  oxidation:  IV.  Zirconium and 114,  hafnium carbides’,  J. Electrochem. Soc,  1967,  (10), 1030-1033.  97. S. Shimada and T. Ishil: ‘Oxidation kinetics of zirconium carbide at relatively low temperatures’, J. Am. Ceram. Soc, 1990, 73, (10),  2804-2808.  98. G.  A.  Rama  Rao  and  V.  Venugopal:  ‘Kinetics  and  mechanism of 206,  the oxidation of ZrC’,  J. Alloys Compd,  1994,  (2), 237-242.  99. K. Tamura, T. Ogawa and K. Fukuda: ‘The oxidation behavior of  ZrC coating and powder studied by laser Raman spectroscopy and X-ray diffraction’, J. Nucl. Mater, 1990, 175,  (3), 266-269.  100. A. C. Fox and T. W. Clyne:  ‘Oxygen transport by gas permeation  through the zirconia layer in plasma sprayed thermal coatings’, Surf. Coat. Technol., 2004, 184, (2-3), 311-321.  barrier  101. S. N. Karlsdottir and J. W. Halloran:  ‘Formation of oxide scales  on  zirconium  diboride-silicon  carbide  composites  during  oxidation: relation of subscale recession to liquid oxide flow’, J. Am. Ceram. Soc, 2008, 91, (11), 3652-3658.  102. E. Zapata-Solvas, D. D. Jayaseelan, P. M. Brown and W. E. Lee:  ‘Effect of La2O3 addition on long-term oxidation kinetics of ZrB2 SiC and HfB2-SiC ultra-high Ceram. Soc, 2014, 34,  temperature  ceramics’,  J. Eur.  (15), 3535-3548.  103. W. E. Lee, E. Giorgi, R. Harrison, A. Maitre and O. Rapaud:  ‘Ultra-high  temperature  ceramics:  materials  for  extreme  environment  applications’  ‘Nuclear  applications  for  ultra-high  temperature  ceramics  and MAX  phases’,  391-415;  2014,  Hoboken, NJ, John Wiley & Sons Inc.  Ha r r ison and Lee  P rocess ing and p rope r t ies o f Z rC , Z rN and Z rCN ce ram ics  3 0 7  Advances in App l ied Ce ram ics  2016  VOL 115  NO 5  \\x0c']"
},{
  "_id": 219,
  "PDF": "Processing, Mechanical Properties and Oxidation Behavior of TaC and HfC Composites Containing 15 vol_ TaSi2 or MoSi2.pdf",
  "Text": "['See discussions, stats, and author profiles for this publication at : https://www .researchgate.net/publication/250011712  Processing, Mechanical Properties and Oxidation Behavior of TaC and HfC  Composites Containing 15 vol% TaSi2 or MoSi2  Article · June 2009  DOI: 10.1557/  jmr.2009.0232  CITATIONS  79  5 author  s, including:  Diletta Sciti  Italian National Research Council  204   PUBLICATIONS\\xa0\\xa0\\xa04,100   CITATIONS\\xa0\\xa0\\xa0  SEE PROFILE  Stefano Guicciardi  Italian National Research Council  123   PUBLICATIONS\\xa0\\xa0\\xa02,905   CITATIONS\\xa0\\xa0\\xa0  SEE PROFILE  Some of the authors of this publication are also working on these related projects:  C3HARME View project  ST2G Solar Thermionic Thermoelectric Gener ator View project  READS  260  Laura Silvestroni  Italian National Research Council  110   PUBLICATIONS\\xa0\\xa0\\xa02,026   CITATIONS\\xa0\\xa0\\xa0  SEE PROFILE  A. Bellosi  Italian National Research Council  234   PUBLICATIONS\\xa0\\xa0\\xa05,634   CITATIONS\\xa0\\xa0\\xa0  SEE PROFILE  All content following this page was uploaded by Stefano Guicciardi on 20 April 2015.  The user has requested enhancement of the downloaded file.  \\x0c', 'Processing, mechanical properties and oxidation behavior of TaC  and HfC composites containing 15 vol% TaSi2 or MoSi2  Diletta Sciti,a) Laura Silvestroni, Stefano Guicciardi, Daniele Dalle Fabbriche, and Alida Bellosi  CNR-ISTEC, Institute of Science and Technology for Ceramics, I-48018 Faenza, Italy  (Received 22 October 2008; accepted 23 February 2009)  Fully dense HfC and TaC-based composites containing 15 vol% TaSi2 or MoSi2 were  produced by hot pressing at 1750-1900  \\x0e  C. TaSi2 enhanced the sinterability of the  composites and nearly fully dense materials were obtained at lower temperatures than in  the case of MoSi2-containing ones. The TaC-based composites performed better than HfC  composites at room temperature, showing values of mechanical strength up to 900 MPa and a fracture toughness of 4.7 MPa\\x01m1/2. However, preliminary oxidation tests carried  out in air at 1600  \\x0e  C revealed that HfC-based composites have a superior high  temperature stability compared to TaC-based materials.  I.  INTRODUCTION  Tantalum and hafnium carbide are refractory transition  metal carbides from the fifth and the fourth group of the periodic table. Because of their melting point (>3900  \\x0e  C),  high hardness, elastic modulus, and resistance to chemical attack,1-3  they are  considered candidates  for  extremely  high-temperature applications such as rocket nozzles and  scramjet components, where the operating temperatures  can be in excess of 3000  \\x0e  C. Monolithic TaC or HfC have  been traditionally difficult  to densify due  to their high  covalent bonds and low self-diffusion coefficient. There fore, hot pressing has been typically used to enhance the  densification of  these compounds. Several studies report  the  results of  attempts  to improve  the densification of  carbides since the 1960s. TaC was densified up to 90%-  95% by hot pressing, at temperatures in the range 2200- C.4-6 Other studies7,8  3000  \\x0e  report  that  the temperature  needed  for  densification was  strongly  reduced  (1400-  1800  \\x0e  C) with small  amounts of Fe, Mn, Co,  and Ni.  However, the presence of the liquid phase formed by these  additives may  cause  exaggerated  grain  growth, which  leads to entrapped porosity within large grains in the final  stage of densification and prevents the ceramics from reaching full density.6 Furthermore, the presence of metal lic phases at  the grain boundaries is known to limit  the  usefulness of ceramics for high-temperature applications.  Recently, the addition of C and B4C to enhance the densification of TaC has been reported in the literature.9 These  additives are believed to favor the elimination of the sur face oxides present on the starting particles. As a matter of  fact,  the onset  temperature for densification of TaC was  lowered through the use of either C or B4C additives, or  both. The maximum density with 0.36 wt% B4C addition  was 98% at 2200  \\x0e  C, accompanied by rapid grain growth.  Similar densification problems were encountered for haf nium carbide. Opeka et al. sintered HfCx compounds by hot  pressing at 2500 sizes of the order of \\x1840-60 mm.10,11 Lack of control of the C, achieving materials with mean grain  \\x0e  microstructure, due to grain coarsening and entrapped po rosity,  is thus the main issue for  these compounds. HfC based composites have been recently sintered to full density  either by spark plasma sintering or pressureless sintering,  thanks  to  additions  of MoSi2  as  the  sintering  agent.  Amounts of between 3 and 20 vol% of molybdenum sili cide allow densification to occur at 1900-1950  \\x0e  C achiev ing fine microstructures with grain size of the order of few micrometers (1-3 mm).12,13 Due to the fact  that  these car bides are mostly used for military applications, published  values for mechanical properties are scarce, especially for  TaC-based materials. Moreover,  the  reported values  are  widely  scattered  because  of  differences  in  preparation  method, mean grain size, and final density. The strength of TaC reported in the 1960s by Santoro14 was obtained on  specimens prepared by carburization of Ta metal. Santoro  studied the variation of  the TaCx properties with carbon  content and found that strength and hardness follow oppo site trends. A minimum flexure strength of 689 MPa was  measured for TaC0.8 corresponding to a maximum hardness of \\x1829 GPa. Recent hardness and fracture toughness data are 21-27 GPa and 3.5 MPa\\x01m1/2, respectively. Measured on vacuum plasma sprayed TaC reported by Balani et al.15 values for Young’s modulus vary from 294 GPa16 to 550 GPa.17 As  far  as  concerns monolithic HfCx  composites,  strength values and elastic modulus reported by Opeka and  coworkers were around 150 MPa and 380 GPa, tively.10 Most  respec likely the strength was negatively affected  by the coarse microstructure. Values of mechanical strength  of 450-465 MPa were found for pressureless sintered HfC  a)Address all correspondence to this author.  e-mail: diletta.sciti@istec.cnr.it  DOI: 10.1557/JMR.2009.0232  J. Mater. Res., Vol. 24, No. 6, Jun 2009  © 2009 Materials Research Society  2056  \\x0c', 'composites  containing  5-10  vol% MoSi2.12 The these composites was in the range 415-434  elastic  modulus  for  GPa  and the found to be 3.5-3.6 MPa\\x01m1/2.12 As fracture toughness measured by the CNB  method was  for  the  mechanical properties, available studies on oxidation are  still scarce in the literature, especially for TaC. It is known  that TaC oxidizes to form Ta2O5 which has a relatively low C. Wang et al.18 studied the high  melting point at 1800  \\x0e  temperature oxidation of HfC-TaC composites. According  to their study, at T < 1800 C the outermost oxide layer was constituted of Ta2O5, while at T > 1800 Ta2Hf6O19 and HfO2 was found.  \\x0e  \\x0e  C a mixture of  More information is available on HfC. Hafnium car bide oxidizes  to monoclinic HfO2 which is one of  the  most stable refractory compounds, with a melting point  of 2900  \\x0e  C. The high resistance to oxidation of mono lithic HfC was proved either by conventional oxidation tests carried out between 1400 and 2060 C19 or by arc C.20 Similarly, HfC  \\x0e  jet  tests between 2400 and 2700  \\x0e  composites  containing 5 vol% MoSi2  showed a good  thermal stability when tested in an arc jet facility up to C.21 Finally, the addition of TaC to HfC materials  2400  \\x0e  indicates a general trend toward enhanced oxidation and spalling with increasing TaC content.22 Other  surface  studies confirm that  the addition of TaC to other ultra refractory compounds such as ZrB2 or ZrB2-SiC generally leads to a worsening of the oxidation resistance.23  In this work,  the densification, microstructure, and me chanical properties of TaC and HfC-based composites con taining  15  vol% MoSi2  or  TaSi2  are  presented  and  discussed. Preliminary oxidation studies are also conducted  in air at 1600  \\x0e  C in a bottom-up furnace to test and compare  the high temperature stability of these materials.  II. EXPERIMENTAL PROCEDURE  Commercial powders were used to prepare the ceram ic composites: cubic HfC (Cerac Inc., Milwaukee, WI), particle size range 0.2-1.5 mm; cubic TaC (Cerac Inc., Milwaukee, WI), particle size range 0.2-1.5 mm; hexagonal  TaSi2 (ABCR, GmbH & Co, Karlsruhe, Germany), particle size <45 mm; (<2 mm, Aldrich, tetragonal MoSi2 Steinbeim, Germany), particle size range 0.3-5 mm and oxygen content \\x181 wt%. The  following compositions  were prepared (vol%): TCM : TaC þ 15 MoSi2  ;  TCT : TaC þ 15 TaSi2  ;  HCM : HfC þ 15 MoSi2  ;  HCT : HfC þ 15 TaSi2  :  The powder mixtures were ball milled for 24 h in  absolute  ethanol using zirconia milling media. Subse quently,  the mixtures were dried in a rotary evaporator  and sieved through a 60-mesh screen. Hot-pressing was conducted in low vacuum (\\x18100 Pa) using an induction heated graphite die with a constant uniaxial pressure of  30 MPa, heating rate 20  \\x0e  C/min, and free cooling. For  each composition,  the maximum sintering temperature  was set on the basis of the shrinkage curve, as explained  later. The bulk densities were measured by Archimedes’  method. Crystalline phases were identified by x-ray dif fraction (Siemens D500, Germany). The microstructures  were  analyzed  using  scanning  electron microscopy  (SEM, Cambridge S360, Cambridge, UK)  and energy  dispersive spectroscopy (EDS, INCA Energy 300, Oxford  Instruments, UK)  on  polished  surfaces. Mean  grain  sizes,  amount  of  porosity,  and  amount  of  secondary  phases were  determined  through  image  analysis  on  SEM micrographs of polished surfaces using a commer cial software program (Image Pro-Plus 4.5.1, Media Cy bernetics, Silver Spring, MD). At  least 100 grains per  specimen were measured for  the determination of  the  mean grain size.  Vickers microhardness (HV1.0) was measured with a  load  of  9.81 N,  using  a Zwick  3212  tester. Young’s  modulus (E) was measured by the resonance frequency method on 28 \\x02 8 \\x02 0.8 mm specimens using an HP gain-phase analyzer. Fracture toughness (KIc) was evalu ated using chevron-notched beams (CNB) in flexure. The test bars, 25 \\x02 2 \\x02 2.5 mm (length by width by thick ness,  respectively), were notched with a 0.1 mm-thick  diamond saw;  the chevron-notch tip depth and average  side length were about 0.12 and 0.80 of the bar thickness,  respectively. The specimens were fractured using a semi articulated silicon carbide four-point fixture with a lower  span  of  20 mm and  an  upper  span  of  10 mm using  a  screw-driven  load  frame  (Instron mod.  6025, High  Wycombe, UK). The specimens,  three for each compos ite, were loaded with a crosshead speed of 0.05 mm/min. The “slice model” equation of Munz et al.24 was used to  calculate KIc. On the same machine and with the same the flexural strength (s) was measured at C on chamfered bars 25 \\x02 2.5 \\x02 room temperature and at 1200 2 mm (length \\x02 width \\x02 thickness, respectively), using a  fixture,  \\x0e  crosshead speed of 0.5 mm/min. The high-temperature  strength was tested at 1200  \\x0e  C under flowing argon pro tective gas. Before the bending test, a soaking time of  18 min was set  to reach thermal equilibrium. For each  composition, five specimens were tested at room temper ature and three specimens at 1200  \\x0e  C.  13 \\x02 2.5 \\x02 2 mm bars in static air, exposure of 15 min in a The high temperature stability was tested at 1600 C on  \\x0e  bottom-up loading furnace box. All  the specimens were  previously cleaned in acetone. The mass of the specimens  was measured before and after exposure. The microstruc tural modifications  induced  in  the  oxidized  specimens  were evaluated by SEM-EDS on the surface and cross  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2057  \\x0c', 'section. The specimens were located in the furnace when  the maximum temperature was reached and then removed  and air quenched after an exposure time of 15 min. The  degradation after oxidation was evaluated by comparing  the pristine strength of  the as-sintered specimens and the  retained strength of the oxidized specimens. The specimens  were  tested at  room temperature  in three point-bending  tests using as-sintered specimens with the same dimensions  as the oxidized ones.  III. RESULTS AND DISCUSSION  A. Densification  Density data and sintering cycles are summarized in  Table I. Shrinkage curves versus  time are displayed in  Figs. 1(a)-1(d). TaC-MoSi2 (TCM) started the densifica tion at 1690 C [Fig. 1(a)] and needed a temperature of C to achieve a density of 12.8 g/cm3. TaC-TaSi2 composite (TCT) started shrinking at 1400 C [Fig. 1(b)]  \\x0e  1850  \\x0e  \\x0e  and  required  a maximum temperature of 1750 g/cm3. HfC-MoSi2 C [Fig. 1(c)] and required a  \\x0e  C to  achieve  a  density  of  13.3  (HCM)  started shrinking at 1760  \\x0e  maximum temperature of 1900 C to achieve a bulk density of 11.7 g/cm3. Finally, HfC-TaSi2 (HCT) started  \\x0e  shrinking at 1600 C [Fig. 1(d)] and completed the densiC with a final bulk density of 12.0 g/cm3.  \\x0e  fication at 1760  \\x0e  The relative densities, determined as  the ratio between  bulk and theoretical densities calculated on the basis of starting compositions, were >96% for all the composites  (Table I). However,  since post-sintering SEM analyses  ascertained the presence of extra phases (SiC and SiO2)  having much lower density than the starting ones,  the  relative  densities  indicated in Table  I  can  be  signifi cantly underestimated, as discussed in the following sec tion. The densification curves show that  the addition of  TaSi2  is more effective for  sintering of  these carbides  compared to MoSi2, as it  lowers the maximum sintering  temperature of about 100  \\x0e  C for TaC and 140  \\x0e  C for  HfC-based composites. Moreover, the shrinkage of com positions containing TaSi2 started at a much lower tem perature than those containing MoSi2.  B. X-ray diffraction and microstructural features  Microstructural details of the studied compositions are  summarized in Table I.  TaC-MoSi2  (TCM):  The  x-ray  diffraction  pattern  shows  crystalline TaC and MoSi2  after  sintering. The  fracture and polished sections [Figs. 2(a) and 2(b)] show  TABLE I. Compositions, sintering parameters, density data, and microstructural parameters.  Label  Matrix  Additive  Sintering cycle  (  \\x0e  C/min)  Ton (  \\x0e  C)  Experimental density  (g/cm3)  Relative  density (%)  Matrix m.g.s. (mm) \\x181.2 \\x182.5 \\x181.2 \\x180.8  Additive m.g.s. (mm) \\x180.9 \\x181.9 \\x181.4 \\x180.2  TCM  TaC  MoSi2  1850/3  1690  12.8  96.3  TCT  TaC  TaSi2  1750/9  1400  13.3  97.3  HCM  HfC  MoSi2  1900/10  1760  11.7  99.9  HCT  HfC  TaSi2  1760/10  1600  12.0  98.6  Ton = temperature at which the shrinking starts; m.g.s. = mean grain size.  FIG. 1. Sintering curves for TaCand HfC-based materials. (a) TaC-MoSi2, (b) TaC-TaSi2, (c) HfC-MoSi2, and (d) HfC-TaSi2.  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2058  \\x0c', 'D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  FIG. 2. TaC-MoSi2 composite:  (a)  fracture and (b) polished section.  Inset: SiC formation in the proximity of a MoSi2  inclusion. TaC-TaSi2  composite: (c) fracture and (d) polished section. Dark contrasting features are SiC or SiO2 phases. Inset: EDS spectrum collected from the TaSi2  phase.  that this composite had a fine and homogeneous microstructure. The TaC mean grain size was around 1 mm  while the MoSi2 particles (dark contrasting phase) had a mean grain size of about 1 mm and a grain size range of 0.1-4.8 mm. Despite the low value of  relative density,  an irregular shape. TaSi2 is known to a have a brittle to ductile transition similar to MoSi2.25 This feature allowed the silicide to accommodate in the voids left by the matrix  grains. According to image analysis, the volumetric amount  of TaSi2 in the final microstructure was around 12 vol%,  little or no porosity was detected by SEM. The discrep that  is,  slightly lower  than the initial composition, 15%.  ancy between SEM observation and the relative density  Dark contrasting features were identified as SiC and SiO2  value reported in Table I is due to a considerable amount  phases. Considering the volumetric amount of these phases  of silica pockets in the microstructure, about 2-3 vol%,  as verified by image analysis. SiO2 contamination has  already been observed in all the composites containing secondary phase.26 Large pockets of SiC  as  MoSi2  a  were also detected by SEM-EDS analysis  [see inset  Fig. 2(b)]. Considering the amount of SiC and SiO2  in  in  the calculation of the theoretical density (0.5 vol% SiC,  2 vol% SiO2), the final relative density increases to 99%,  which is more consistent with the absence of porosity  observed by SEM analysis.  TaC-TaSi2 (TCT): TaC and TaSi2 were the only crys talline phases detected after sintering in the x-ray diffrac (2 vol% SiC, 0.5 vol% SiO2) in the calculation of the theotheoretical density is \\x1813 g/cm3  retical density,  the final  and hence the relative density is 99%. This value of relative  density is more consistent with the low quantity of porosity  ascertained by SEM observations.  HfC-MoSi2  (HCM): Reflections from crystalline HfC  and MoSi2 were identified in the x-ray diffraction pattern  of  this composite. A fine-grained microstructure with no  porosity was detected on fracture and in polished sections [Figs. 3(a) and 3(b)]. The mean grain size of HfC was 1.2 mm.  The MoSi2 phase showed a tendency to form large pockets with dimensions up to 5-6 mm. At  inter the  tion spectrum. Fracture and polished surfaces, displayed in  face  between  the  silicide  and  the matrix, SiC grains  Figs. 2(c) and 2(d), put in evidence that little or no porosity  formed. Different  from TCM and TCT,  no  residual  was present  in the microstructure. The fracture surface larger grains (5-7 mm) were transgranularly showed that fractured, while smaller grains (1-3 mm) were intergranu larly fractured. The mean grain size of the carbide phase was 2.5 mm, but coalescence of grains led to the formation of larger grains (up to 6-7 mm). TaSi2 particles are recognizable as having a slightly darker contrast compared to  TaC grains and could form large pockets as wide as 3- 8 mm [Fig. 2(d)]. It can be noted that these particles have  silica was generally detected in this material.  HfC-TaSi2  (HCT): According  to  x-ray  diffraction  (Fig. 4), no reflections  from TaSi2 were detected after  sintering. Moreover, collected reflections do not corre spond to those of pure HfC (PDF No. 39-1491) being  clearly shifted toward higher angles. This feature is displayed in the spectrum collected in the range 2y = 70-  120  C where it was possible to distinguish at  least  two  \\x0e  different peaks in each group of  reflections,  in addition  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2059  \\x0c', 'to uncertain residual reflection from pure HfC. The unit  cell parameters of this newly formed phase were a = 4.623 and 4.420 A˚ , that is, shorter than those of pure HfC (a = 4.637 A˚ ), which indicates a contraction of the unit cell.  The presence of  these additional  reflections was  inter preted as due to the formation of  (Hf, Ta)C solid solu tions. On the basis of the Vegard’s rule for the HfC-TaC  system and hypothesizing that only Ta can enter the HfC  structure, corporated into HfC grains was \\x1810 at.% and \\x1820 at.%, it can be estimated that the amount of Ta in giving  (Hf0.9Ta0.1)C and  (Hf0.8Ta0.2)C as  the  compo sitions of  the two identified solid solutions. The micro structure  of  this  composite  is  constituted  of  squared  carbide grains, dark contrasting phases (i.e.,  low density  phases) at  triple junctions, and residual HfO2 particles.  [Figs. 5(a)-5(c)]. The matrix grains have a mean grain size <1 mm (Table I),  that  is,  lower than the HCM com posite (Table I). This refinement of the microstructure is  certainly  related  to  the  lower  sintering  temperature  needed to complete  the densification (1760  \\x0e  C versus  1900  \\x0e  C). The formation of solid solutions was investi gated  by EDS analysis. Although  there  is  substantial  overlapping of hafnium and tantalum peaks in the ener gy spectrum, Ta was detected in most of  the  carbide  grains [Fig. 5(d)]. Quantitative analyses confirmed that  the Ta  content  can vary between 6 and 15 at.%. The  identification of TaSi2 particles was difficult  in this mi crostructure, due to the very similar contrast displayed  by HfC and TaSi2 and overlapping of EDS signals col lected from Hf, Ta, and Si around 1.7 keV. However, the  micrographs  show that  the  large TaSi2  inclusions ob served in TCT are completely absent  in HCT,  in agree ment with  the  findings  of  x-ray  diffraction. Residual  secondary  phases  observed  at  triple  junctions  are Si based phases that locally formed large pockets, as shown  in Fig. 5(e). When these regions were analyzed at  low  accelerating voltage (EHT = 5 keV), only the peaks of  Si and C could be detected [see the inset  in Fig. 5(e)].  On the other hand, when the  spectrum was  collected  at higher  accelerating voltage  (EHT = 20 keV),  addi tional peaks  from Hf were also identified. This makes  it difficult  to ascertain whether  the C peak is collected  from the  underlying  carbide matrix  or  not,  that  is,  whether the Si-based phase is silicon or SiC. As carbon  FIG. 4. X-ray diffraction pattern for  the HCT composite in the 2-y range 70  \\x0e  -120  \\x0e  . The underlying spectrum refers to pure HfC according to  PDF No. 39-1491.  FIG. 3. HfC-MoSi2 composite:  (a)  fracture and (b) polished section.  SiC formation is observed at  the interface between MoSi2 and HfC.  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2060  \\x0c', 'sources are  available in the present  system,  it  is  like ly  that  a  partial  or  complete  carburization  of  silicon  occurs.  C. Sintering mechanisms  Densification results  show that  the  addition of  sili cides  strongly improves  the densification of TaC and  HfC carbides. The  sintering temperatures  are between  1750 and 1900 C, that is, they are much lower those reported in the literature (>2200  \\x0e  than  \\x0e  C4,6,10,11). The  shrinkage  curves of TCM, HCM,  and TCT display a  slope change, that is, an acceleration in the sintering rate  [Figs. 1(a)-1(c)]. This is usually attributed to the forma tion  of  liquid  phase, which  activates matter  transfer  tion is completed at T \\x19 1740 mechanisms. The slope change indicates that liquid formaC for HCM, T \\x19 1700 and T \\x19 1500  \\x0e  \\x0e  C  for TCM,  \\x0e  C for TCT. HCT does not  display any transition and the sintering rate,  that  is,  the  curve slope,  is higher than the other systems [Fig. 1(d)].  Microstructural  features of  the compositions under  in vestigation confirm the formation of Sior SiO2-based  phases, which are presumed to be liquid at  the sintering  temperatures. All these observations lead to the conclusion  that  densification  of  these  carbides was  activated  by  liquid phase mechanisms.  In one case, HCT,  formation  of solid solution was also observed. No solid solutions  were instead detected for HCM, TCM, and TCT, at least  at the resolution of the used techniques. In the following,  densification mechanisms are hypothesized for  the four  systems.  TCM: The inhibition of sintering in nonoxide ceramics  such as carbides is generally attributed to the presence of oxide impurities on the powder particle surface.9 The pres ence of significant amounts of silica in the microstructure  suggests  that oxidation of MoSi2 particles can help the  removal of TaC particles  surface oxides, promoting the  densification. Silica  could also favor  the particle  rear rangement, acting as grain lubricant during the initial stage  of the densification. However,  the morphology of MoSi2  particles displaying very low dihedral angles suggests that  the silicide can form a liquid phase at the sintering temper ature. It has already been hypothesized that interaction of MoSi2 with carbon can induce liquid phase formation.27,28 Carbon is certainly present as an impurity in the starting  carbide powders (usually produced through carbothermal  reduction of the metal oxide) and is abundant in the graph ite furnace (die and rams). According to the calculated C-MoSi2 pseudo-binary phase diagram29 above 1700 a liquid phase forms at any C/MoSi2 concentration.  \\x0e  C,  If  MoSi2  is more  abundant  than C,  a mixture  containing  liquid phase + MoSi2 + SiC is predicted. The formation  of SiC in close proximity to the MoSi2 phase  and the  temperature of  liquid phase formation extrapolated from  the shrinkage curve are consistent with predictions of the  phase diagram.  TCT: The presence of SiO2 pockets in the microstruc ture suggests that TaSi2 reacts with the surface oxides of  FIG. 5. HfC-TaSi2 composite:  (a)  fracture and (b) polished section showing squared carbide grains and secondary low density phases at  triple  junctions indicated by arrows, (c) HfO2 impurities, (d) a typical EDS spectrum collected on HfC grains revealing the presence of Ta, (e) a SiC  phase deriving from carburization of silicon and relative EDS spectrum collected at  low energy.  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2061  \\x0c', 'TaC, cleaning the powder particle surface from oxygen  impurities, as concluded for TCM. However,  the earlier  start  of  densification  (1400  \\x0e  C)  indicates  that  further  mechanisms may be active. In a previous study concerning boride-TaSi2 composites,30 can undergo decomposition reacting with CO typically  it was found that TaSi2  present  in graphite-based furnaces.  In the present case,  the following reaction is hypothesized to occur: TaSi2 þ COðgÞ ¼ TaC þ Sið1Þ þ SiOðgÞ favorable T > 1100 oC  :  ð1Þ  The formation of silicon, which is liquid at 1415  \\x0e  C,  is  very close to the onset of densification for this compos ite. The decrease of TaSi2 content after sintering and the  presence of Si-based phases  in the final microstructure  are consistent with the occurrence of reaction (1). Final ly,  the formation of SiC phases could derive from the  carburization of silicon.  HCM: For this system,  it  is hypothesized that densifi cation mechanisms were similar  to those of TCM. Liq uid phase formation around 1740  \\x0e  C and SiC formation  are consistent with C-MoSi2 phase diagram predictions.  Furthermore, previous studies carried out on HfC-MoSi2  composites  evidenced  the  presence  of  Hf-Mo-Si-C  mixed phases, which supports the hypothesis that MoSi2 acts as a medium for matter transfer mechanisms.12  HCT: The absence of TaSi2 after sintering [as verified  by x-ray diffraction (Fig. 4)],  the presence of  silicon based phases, and the enhanced densification behavior  of  this composite suggest  that decomposition of TaSi2  according to reaction (1)  is  active  in this material  as  well. However, TaC derived from reaction (1) was not  detected in the microstructure.  Instead,  (Ta, Hf)C solid  solutions were identified by x-ray diffraction and SEM EDS analysis. This  suggests  that  reaction  (1)  can  be  rewritten for this composite as: TaSi2 þ HfC þ COðgÞ ¼ ðTa; Hf ÞC þ Sið1Þ þ SiOðgÞ  :  ð2Þ  As for TCT,  the SiC phase can form through carburiza tion of silicon. It  is worth noting that  in HCT,  the solid  solution formation leads  to complete decomposition of  TaSi2, different from TCT. It  is likely that solid solution  formation provides an extra driving force that acceler ates the decomposition of TaSi2 and hence the sintering  rate of  this  system. Future TEM studies  are meant  to  confirm these hypotheses.  D. Mechanical properties  The values of  the mechanical properties are summa rized in Table II. Despite the hardness of MoSi2 (12 GPa),31 which is lower than that of TaSi2 (15.6 GPa),32 the hardness of the TaC composites (14.5 GPa) does not  change significantly upon addition of MoSi2 or TaSi2. In  TCM,  the  lower  hardness  of MoSi2  could have  been  compensated for by the  finer matrix mean grain size  (Table I). Also the Young’s modulus of posites (\\x18490 GPa) does not change significantly changthe TaC com ing the silicide phase, despite the differences in the stiffness of MoSi2 (E = 425 GPa)33 and TaSi2 (E = 360 GPa).34 Significant contamination of silica (2-3 vol%)  in TCM could have affected the overall value of  stiff ness. The  fracture  toughness  is  also unchanged when  MoSi2 is replaced by TaSi2,  indicating that  this property  is mainly dictated by the matrix toughness. A relevant  difference is instead observed for the flexural strength at  room temperature (TCM: 900 MPa, TCT: 679 MPa).  It  is likely that  this property was positively influenced by  the refined and more homogenous grain size of the TCM  composite compared to TCT. For  the high temperature  strength test, the TaC-based composites were preliminar ily subjected to a conventional  test  in air at 1200  \\x0e  C.  However, despite the short permanence in temperature,  the  tested specimens were heavily oxidized. To over come this problem, during the subsequent high tempera ture bending tests, the furnace chamber was flushed with  argon to minimize the interaction of  the samples with  oxygen. Nonetheless,  the  high  temperature  flexural  strength was highly reduced compared to the room temis, \\x0036% for TCT and \\x0040% for  perature value,  that  TCM, see Table II.  A direct comparison of properties between HCT and  HCM is not feasible because of solid solution formation  in HCT, whose properties are not known. Nevertheless,  irrespective of the microstructural differences,  the com posites have  close values of  the properties. HCM has  slightly  higher  values  of  (3.80 MPa\\x01m1/2 versus 3.58 hardness (19.6 GPa versus  18.4 GPa) MPa\\x01m1/2) but and toughness  lower  stiffness  (451 GPa)  and strength  (417 MPa)  than HCT (489 GPa and 464 MPa,  respec tively). As  for TaC composites,  the  flexural  strength  tested at 1200  \\x0e  C in argon is  reduced compared to the  room temperature value even if to a lesser extent than in the TaC-based composites: \\x0010% for HCT and \\x0030%  for HCM (Table II).  Generally  speaking,  HfC-based  composites  have  higher  values  of  hardness,  but  lower  toughness  and  TABLE II. Mechanical properties.  Label  HV (GPa) 14.5 \\x06 0.3 14.6 \\x06 0.4 19.6 \\x06 0.5 18.4 \\x06 0.9  E (GPa) 490 \\x06 5 486 \\x06 5 451 \\x06 5 489 \\x06 5  KIc  (MPa\\x01m1/2) 4.7 \\x06 0.1 4.7 \\x06 0.1 3.80 \\x06 0.03 3.58 \\x06 0.03  s  RT (MPa) 900 \\x06 33 679 \\x06 18 417 \\x06 38 464 \\x06 95  s1200  (MPa) 537 \\x06 45 429 \\x06 35 294 \\x06 39 394 \\x06 27  TCM  TCT  HCM  HCT  HV: Vickers hardness; E: Young’s modulus; KIc: fracture toughness, CNB; sRT: room temperature; s1200:  four-point  flexural  strength at  four-point  flexural strength at 1200  \\x0e  C in Ar flux.  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2062  \\x0c', 'strength,  than TaC-based materials. However,  the latter  system shows a higher decrease rate of strength with the  test temperature.  E. Oxidation tests  The oxidation data, composed of weight gain, crystal line phases, and mechanical strength, are summarized in  Table III. Weight gain data indicate that TaC-based com posites were more strongly oxidized than HfC-based ones.  The  change  of  additive  from TaSi2  to MoSi2  had  no  relevant  effect  on  the  overall  oxidation  resistance  of  TaC composites, whose  behavior was mainly  driven  by the characteristics of  the matrix. According to x-ray  diffraction,  the only crystalline phase detected on the sur face of both materials  is Ta2O5, but  in the case of  the  composite  containing MoSi2,  the  crystalline Ta2O5  on  the oxidized surface presented a preferred 00l orientation.  SEM analysis  of  the  fractured  cross  sections  show  morphology and thickness of the oxide formed on TaC based ceramics  [Figs. 6(a)  and 6(b)]. On the  surface,  Ta2O5  crystals  are  embedded in a  silica-based glassy  phase  that derives  from oxidation of MoSi2 or TaSi2  [see insets in Figs. 6(a) and 6(b)]. The extensive oxida tion resulted in a high amount of  low viscosity silicatic  phase, where phase separation occurred, giving rise to  the  formation of  a peculiar  configuration of Ta-oxide  crystals. Ta2O5 crystals grew with a preferred orientation  (particularly in the oxidized surface of TCM) favored by  the immiscibility of this transition metal oxide with glass silicates.35 The porosity amount and size in the oxide  layer appeared to increase from the bulk/oxide interface  to the surface, particularly in the TCM sample, probably  due to the release of gaseous Mo-oxides  in the upper  oxide  layer which is  in contact with the  atmosphere.  The cross section of  the two composites shows that  the  porous scale formed after 15 min of treatment at 1600  \\x0e  C  is 1 mm or even 2 mm thick. It is also apparent, from the  sample cross sections,  that a large volume increase oc curred for sions increase from 2 \\x02 2.5 mm2 to \\x183.7 \\x02 4 mm2). In these materials (TCM cross section dimen fact,  the formation of Ta2O5 is associated with a high the Pilling and Bedworth ratio (R = 2.5)36 that  value of  induces  stresses  in  the  scale. According  to  previous  FIG. 6. Cross sections of  the composites after oxidation.  (a) TCM,  (b) TCT (insets: porous microstructure and surface morphology showing  Ta2O5 crystals embedded in SiO2), (c) HCM, (d) HCT (inset: surface morphology).  TABLE III. Oxidation tests.  Label  Weight  gain/unit  surface (mg/cm2)  Crystalline  phases by XRD  3-pt s as sintered  (MPa) 1091 \\x06 126 739 \\x06 159 533 \\x06 32 564 \\x06 6  3-pt s  after  oxidation  (MPa) 545 \\x06 255  TCM  77.0  Ta2O5  TCT  77.4  Ta2O5  525 \\x06 29 323 \\x06 74  HCM  5.6  HfO2  HCT  15.7  Hf6Ta2O17, HfO2  3-pt s: three point flexural strength.  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2063  \\x0c', 'results on the oxidation of different carbides,37 the oxi dation of TaC could be controlled by an interfacial reac tion that  results in Ta-oxide crystallites bonded directly  to the TaC lattice. The good adherence of  the oxide to  the unreacted core does not allow stress  relaxation by  spalling, but  leads to the opening of the cube edges and  the formation of a Maltese cross. This phenomenon has  been reported for the oxidation of either TaC-based materials36 or Hf-based ceramics.38 This particular oxide mor phology was observed in the TCM composite [Fig. 6(a)],  but not  in TCT. This could be due to the higher amount  of porosity and microcracks  in the latter  [see  inset  in  Fig. 6(b)]  that  favored stress relaxation associated with  the phase reaction from TaC to Ta2O5.  As far as concerns HfC-based composites,  the effects  of  thermal  treatment at 1600  \\x0e  C were milder  than for  TaC materials. No dramatic volume  changes occurred  and the weight gains were much lower  (Table III). The  weight  gain was  higher  for  the  composite  containing  TaSi2 than for MoSi2. The cross sections are displayed  in Figs. 6(c) and 6(d), and in agreement with weight gain  values, the thickness of the oxide layer was HCM (70 mm) than for HCT (150 mm).  lower  for  In the case of  HCM, [Fig. 6(c)] only crystalline HfO2 was detected by  x-ray diffraction (Table III). On the oxidized surface of  HCT, HfO2 or Hf6Ta2O17 crystals were detected without  any  trace  of  a  silica  layer  [see  inset  of  Fig.  6(d)]  (Table III). The mixed oxide is likely to be the result of  oxidation of the (Hf, Ta)C solid solutions. Furthermore,  cracking of the oxide occurred at the corners of the HCT  scale [Fig. 6(d)].  In both the composites, no continuous  protective silica layer formed on the surface. This inter esting feature, which deserves deeper  investigation, has  already been found for HfC-MoSi2 composites tested in C.21 One possible explanation  an arc jet facility at 2000  \\x0e  is that  the formation of a continuous silica layer  is hin dered  by  the  emission  of  gaseous CO2  derived  from  oxidation of the carbides. CO2 species may cause active  oxidation of the silicides, that is, release of gaseous SiO.  To quantify the effect of surface damage on the flexur al strength, room temperature flexural strength tests were  carried out on oxidized bars after  thermal  treatment at  1600  \\x0e  C and compared to as-sintered bars. The values of  the retained strength after oxidation (Table III) are con sistent with the analysis of the oxide morphology and the  weight gain values. No reliable value of strength could  be obtained for the TCT composite, due to the specimens  being damaged in the oxidation process. Considering the  absolute values, TCM has the highest value of  retained  strength followed by HCM and HCT. However, when  compared to the pristine  strength it  turns out  that  the  retained strength of TCM after oxidation is about 50%  of the pristine value, while for HCT and HCM this ratio  is about 57% and 98%,  respectively. The decrease rate  after oxidation is  therefore higher  in TCT than in the  HfC-based  composites. Considering  also  the  strength  results presented in Sec. III. D,  it can be affirmed that at  room temperature the TaC-based composites are stronger  than the HfC-based composites, but  they are also more  affected by high temperature degradation.  IV. CONCLUSIONS  TaC and HfC composites plus additions of 15 vol%  TaSi2 or MoSi2 were densified to 98%-99% of  relative  density by hot pressing in the temperature range 1750-  1900  \\x0e  C. The addition of TaSi2  lowered the maximum  sintering temperature of about 100  \\x0e  C for TaC and 140  \\x0e  C  for HfC-based composites compared to MoSi2. This was  attributed to the decomposition of TaSi2 in the presence  of CO(g) during hot pressing with consequent formation  of Si-based liquid phases. Solid solution formation in the  (HfxTay)C system was observed for the HfC-TaSi2 com posite. HfC-based  composites  have  higher  values  of  hardness, but  lower modulus,  toughness,  and strength  than TaC-based materials. A flexural 4.7 MPa\\x01m1/2 was strength of 900  MPa  and  a  fracture  toughness  of  found for TaC-MoSi2 composite. For all  the composi tions,  a notable decrease of  the  flexural  strength was  observed at 1200  \\x0e  C (from 10%-40% of the room tem perature value). Oxidation tests at 1600  \\x0e  C showed that  the HfC-based composites have a higher oxidation resis tance  than TaC-based  composites. 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Balbo and D. Sciti: Spark plasma sintering and hot pressing of  ZrB2-MoSi2 ultra-high temperature A 475, 108 (2008).  ceramics. Mater. Sci. Eng.,  27. D. Sciti, S. Guicciardi, and M. Nygren: Spark plasma sintering  and mechanical behavior of ZrC-based composites. Scr. Mater.  59, 638 (2008).  28. L. Silvestroni and D. Sciti: Microstructure and properties of pressure less sintered ZrC-based materials. J. Mater. Res. 23, 1882 (2008).  29. X.  Fan, K. Kack,  and T.  Ishigawi: Calculated C-MoSi2  and  B-Mo5Si3 pseudo-binary diagrams for the use in advanced materials  processing. Mater. Sci. Eng., A 278, 46 (2000).  30. D. Sciti, L. Silvestroni, G. Celotti, C. Melandri, and S. Guicciardi:  Sintering and mechanical properties of ZrB2-TaSi2 and HfB2TaSi2 ceramic composites. J. Am. Ceram. Soc. 91, 3285 (2008).  31. D. Sciti, S. Guicciardi, and A. Bellosi: Microstructure and prop erties of Si3N4-MoSi2  composites.  J. Ceram. Proc. Res. 3, 87  (2002).  32. G. Schultes, M. Schmitt, D. Goettel, and O. Freitag-Weber: Strain  sensitivity of TiB2, TiSi2, TaSi2 and WSi2 thin films as possible candidates for high temperature strain gauges. Sens. Actuators, A  126, 287 (2006).  33. M. Nakamura, S. Matsumoto, and T. Hirano: Elastic constants  of MoSi2  and WSi2  single-crystals.  J. Mater.  Sci.  25,  3309  (1990).  34. F. Chu, M. Lei, S.A. Maloy,  J.J. Petrovic,  and T.E. Mitchell:  Elastic properties of C40 transition metal disilicides. Acta Mater.  44, 3035 (1996).  35.  I.G. Talmy,  J.A. Zaykoski, and M.M. Opeka: High temperature  chemistry and oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3 and TaSi2. J. Am. Ceram. Soc. 91, 2250 (2008).  36. M. Desmaison-Brut, N. Alexandre, and J. Desmaison: Compari son  of  the  oxidation  behavior  of  two  dense  hot  isostatically  pressed  tantalum carbide  (TaC and Ta2C) materials.  J. Eur.  Ceram. Soc. 17, 1325 (1997).  37. S. Shimada:  Interfacial  reaction on oxidation of  carbides with  formation of carbon. Solid State Ionics 141-142, 99 (2001).  38. D. Mattia, M. Desmaison-Brut, S. Dimovski, Y. Gogotsi,  and  J. Desmaison: Oxidation  behavior  of  an  aluminium nitride hafnium diboride  ceramic  composite.  J. Eur. Ceram. Soc. 25,  1789 (2005).  D. Sciti et al.: Processing, mechanical properties and oxidation behavior of TaC and HfC composites containing 15 vol% TaSi2 or MoSi2  J. Mater. Res., Vol. 24, No. 6, Jun 2009  2065  View publication stats View publication stats  \\x0c']"
},{
  "_id": 220,
  "PDF": "Processing, properties and arc jet oxidation of hafnium diboride silicon carbide ultra high temperature ceramics.pdf",
  "Text": "['ULTRA-HIGH TEMPERATURE CERAMICS  J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5925 - 5937  Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics  M . G A S C H ELORET Corporation, 690 W. Fremont Ave., Sunnyvale, CA 94087, USA  D . E L L E R B Y , E . I R B Y , S . B E C KM A N NASA, Ames Research Center, MS 234-1, Moffett Field, CA 94035, USA  M . G U SM A N ELORET Corporation, 690 W. Fremont Ave., Sunnyvale, CA 94087, USA  S . J O H N S O N NASA, Ames Research Center, MS 234-1, Moffett Field, CA 94035, USA E-mail: Sylvia.M.Johnson@nasa.gov  The processing and properties of HfB2 -20 vol%SiC ultra high temperature ceramics were examined. Dense billets were fabricated by hot-pressing raw powders in a graphite element furnace for 1 h at 2200 C. Specimens were then tested for hardness, mechanical strength, thermal properties and oxidation resistance in a simulated re-entry environment. Thermal conductivity of the current materials was found to be less than previous work had determined while the strength was greater. Oxidation testing of two flat-face models was conducted, at two conditions, for two 10-min durations each. It was concluded that passive oxidation of SiC plays a role in determining the steady-state surface temperatures below 1700 C. Above 1700 C, temperatures are controlled by the properties of a thick HfO2 layer C(cid:2) 2004 Kluwer Academic Publishers and active oxidation of the SiC phase.  1.  Introduction     Ceramic borides, such as hafnium diboride (HfB2 ) and zirconium diboride (ZrB2 ), are members of a family of materials with extremely high melting temperatures which have been referred to as Ultra High Temperature Ceramics (UHTCs). UHTCs constitute a class of promising materials for use in high temperature applications, such as sharp leading edges on future generations of reentry vehicles, because of their high melting points and relatively good oxidation resistance in reentry environments [1]. Because of the extremely high melting temperatures of HfB2 is 3300 of these diboride materials (i.e., melting temperature C as shown in Fig. 1) hot pressing at temperatures >2000 C and at high pressures is required to produce dense materials. The most extensive collection of work conducted on diborides of Hf and Zr was performed by ManLabs in the 1960’s and 70’s under contract with the Air Force. Some of the early work by ManLabs used very high pressure (1-2 GPa) hot pressing to densify the pure diborides [1]. In addition to the pure diborides, ManLabs investigated the influence of a variety of additives, including C and SiC, on the processing, oxidation resistance and thermal shock resistance of Hf and Zr diborides [2]. Their work, in both static furnace experiments and arc jet testing, showed that the addition of SiC improved the     oxidation resistance of Hf and Zr diborides. SiC additions ranging from 5-50 vol% were studied. The work indicated that 20 vol% of a SiC particulate additive resulted in the optimum oxidation resistance [3]. The SiC additions were also found to improve the processing of the materials by reducing the maximum temperature and pressure required to densify the materials and by reducing grain growth of the diboride phase [3]. A number of other studies have been conducted on the oxidation resistance of Hf and Zr diboride materials of various compositions, but these have focused primarily on static or flowing air studies under ambient pressure, and few studies have evaluated their oxidation behavior in simulated reentry environments [4-10]. Static or flowing air oxidation studies under ambient conditions may not provide an accurate representation of a material’s behavior in the unique environment encountered during re-entry. In addition, the relative oxidation resistance of two materials measured in a typical furnace oxidation study may not pertain when testing the same two materials in a reentry environment. Arc jet facilities, on the other hand, provide sustained conditions that are similar to the aerothermal environment experienced during reentry. The testing results are used to understand the thermal performance of materials and systems under controlled aerothermal heating conditions. Arc jet testing is also used to validate  0022-2461  C(cid:2) 2004 Kluwer Academic Publishers  5925  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  properties of the materials have been evaluated and compared to the results from heritage materials of similar compositions (SHARP-B2). Finally, a preliminary investigation of the oxidation behavior of this composition in a simulated reentry environment was begun and testing was conducted at NASA Ames Research Center (ARC) in the Interactive Heating Facility (IHF).  2. Experimental procedure  The raw powders used in this work were −325 mesh HfB2 from Cerac Inc. and 1-2 µm SiC from H.C. Starck Inc. Analysis of the impurities present in the materials was conducted by NSL Analytical Services (Cleveland, OH). Samples were analyzed, using direct current plasma emission spectroscopy to determine metallic impurities, and using inert gas fusion to determine oxygen and carbon impurities. Analysis of the crystalline phases present in the raw and mixed powders was performed with X-ray Diffraction (Scintag X-ray diffractometer) using Cu Kα radiation. Particle size of the raw and milled powders was measured using laser light scattering (Malvern). Light scattering patterns were analyzed with computer software that used the Mie theory. Prior to hot pressing, the HfB2 -20 vol% SiC raw powders were wet-milled with WC milling media in either a planetary mill (Fritsch) or an attrition mill (Union Carbide). The milled powders were carefully dried to prevent phase segregation between the HfB2 (ρ = 11.12 g/cc) and the SiC (ρ = 3.2 g/cc). After drying, the powders were loaded into either 25 or 50 mm diameter graphite dies, lined with graphfoil. Hot pressing cycles were initiated in a graphite element resistance furnace (Thermal Technologies) with vacuum levels of <200 milliTorr. Above 1600 C, the partial pressure of carbon in the furnace increased significantly and degraded the vacuum. Consequently the furnace was back-filled with one atmosphere of inert gas (argon or helium), to preserve the graphite element and insulation. Typical furnace conditions required for densification were 2200 C for 1 h at 25 MPa. The density of the hot pressed billets and test specimens was measured using the Archimedes method and He pycnometry (Micromeritics Accupyc 1330). The phases present in the sintered materials were characterized using XRD. After the samples were crosssectioned and polished to a 1-µm finish, the microstructure of consolidated samples was characterized using optical microscopy and Scanning Electron Microscopy (FEI ESEM 30) with EDX analysis. Mechanical and thermal property test specimens were prepared with diamond tooling, and ground to a final surface finish in accordance with ASTM C1161 (Chand Kare Technical Associates, Worcester, MA). Vickers hardness was determined using indentations created by a Shimadzu HSV-30. Hardness values were calculated from the average length of the diagonal lines across an indentation made using a load of 49 N applied over a 15 s interval. Hardness indents were made in five groups of 3, starting at the outside edge and working towards the center of a sample, to provide a qualitative evaluation of any radial variation in the material microstructure.        Figure 1 Hafnium diboride phase diagram.  the numerical models of materials and systems that are used as design tools [11]. Arc jet testing provides the best ground-based simulation of the reentry environment. Nevertheless, there are a number of differences between the arc jet environment and the reentry environment that must be accounted for when designing an arc jet experiment, and when interpreting the data. For example, surface catalycity can play a more significant role during arc jet testing than it does in flight, because a higher proportion of the air molecules are dissociated in the arc jet environment. NASA Ames began work on UHTCs in the early 1990’s. In 1997 and 2000 Ames demonstrated the use of UHTC sharp leading edges on the Sharp Hypersonic Aero-thermodynamic Research Probe-Ballistic experiments 1 and 2 (SHARP-B1 and SHARP-B2) [12]. The SHARP-B1 vehicle tested a HfB2 -20 vol%SiC nose tip with a 3.5 mm radius. The SHARP-B2 vehicle tested three different leading edge materials, HfB2 -SiC, ZrB2 SiC and ZrB2 -SiC-C. For the sharp wing leading edge applications envisioned for these materials on future reentry vehicles, in addition to oxidation resistance, high thermal conductivity is desirable. In the family of UHTCs, the diborides typically have the highest conductivities. High conductivity improves a material’s thermal shock resistance by reducing temperature gradients and thermal stresses in the material. A higher conductivity also allows more energy to be conducted away from the stagnation point on the wing leading edge. That energy is subsequently reradiated away from lower temperature regions, along the outer surface of the wing leading edge component. This process may enable the vehicle to operate under higher heat flux conditions, improving its performance. The current study has investigated a method of processing HfB2 -20 vol% SiC particulate composites, using conventional hot pressing. Mechanical and thermal  5926  \\x0c', '      quency method, using a GrindoSonic on 3 × 4 × 45 mm Elastic moduli were measured by the resonance frebars. Two methods were employed to measure flexural 1499, using 25.4 mm diameter ×1 mm thick disks) and strength-bi-axial flexure testing (according to ASTM 4-point bend testing (3 × 4 × 45 mm bars according to ASTM C1161), with a crosshead speed of 0.5 mm/min. Estimates of the reliability of the materials were made and represented by a Weibull modulus that was calculated following ASTM C1239. Thermal conductivity of the materials was measured in the range -130 to 2000 C (sample size: 1.27 cm dia. × 0.1 cm thick) using the laser flash technique according to ASTM C201. Coefficient of thermal expansion was measured using a pushrod dilatometer in the temperature range of −130 to 1500 3 × 4 × 10 mm) according to ISO 17562. All thermal C (sample size: property measurements were performed by Netzsch Instruments (Boston, MA). Finally, arc jet testing was performed on machined flat face models, to evaluate the oxidation/ablation behavior of the HfB2 /SiC material. Fig. 2 shows a schematic of the types of arc heaters in use at Ames Research Center. The arc heater produces a hightemperature gas stream by combined radiative, conductive and convective heat transfer from a high-voltage DC electric arc discharge to a gas flowing through a cooled column. The facilities at Ames are capable of input power levels of 20 to 60 MW for up to 30-min durations. The facilities can also test a variety of sample sizes and shapes, such as 80 × 80 cm 2-dimensional articles for panel and seal applications, or 3-dimensional models. The Ames Arc Jet facility can also test samples up to 60 cm in diameter, for aero surface control/response applications. Examples of various HfB2  ULTRA-HIGH TEMPERATURE CERAMICS  UHTC based models that have recently been tested at Ames are shown in Fig. 3. A photo of an as-machined flat face arc-jet model is shown in Fig. 4. The model is 25.4 mm in diameter and the overall height is 8 mm. The notch in the base of the model, shown in the lower left hand picture in Fig. 4, is used to pin the model into the holder. Models were placed in SiC-coated graphite holders (shown in the right hand image of the same figure) which enabled test durations in excess of 10 min; an uncoated graphite holder would have allowed only for a few minutes of testing. Models and holders were then attached to a water-cooled arm (sting), as shown in Fig. 5, that moved the models in and out of the plasma stream. A variety of instrumentation was used to calibrate the arc jet conditions and to measure the thermal response of the materials. The cold wall heat flux values, shown in Table I, were measured using a copper Gardon gage and are referenced to a 76 mm diameter hemisphere. However, the hot wall heat flux values at the model’s surface are different from these calibrations, because of differences in model geometry and differences between the catalycity of the models and the catalycity of the copper Gardon gage. Detailed discussion of the heat flux at the surface of the model is beyond the scope of this work. To differentiate the two conditions, for this text, the first will be designated as “low” and the second will be designated as “high” condition. Two 1-color and one 2-color optical pyrometers were used to make surface temperature measurements during the tests. Based on previous experiments conducted by ManLabs, an emittance of 0.65 was assumed for the 1-color pyrometer [13]. In general there was excellent agreement between the different pyrometers, so  Figure 2 Schematic of the Ames Arc Jet facilities, showing the basic construction of an arc jet. Figure also demonstrates the various nozzle types  available for performing testing on 3-dimensional samples or flat panels [13].  5927  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 3 The family of arc jet models (flat-face, cone and wedge) fabricated and tested at NASA Ames.  Figure 4 Pre-test photo of a flat face arc jet sample, top, side and in a SiC coated graphite holder. Graphite pins were used to hold the sample in place.  To prevent the pins from eroding, the ends were coated with a SiC paste (lighter color spots on the holder).  Figure 5 Photo of a flat-face sample during an arc jet run. Sample and holder are shown attached to a water-cooled sting arm, which serves to move  the sample in and out of the flow.  data from each pyrometer is not shown in every figure. Instrumentation (such as thermocouples or optical pyrometers, for making in-depth or back face temperature measurements) was placed inside a cavity within the sting arm, for protection during the experiments.  For each condition shown in Table I, a model was tested twice for 10-min for a total run time of 20 min. After each exposure, the models were weighed and optical images of their surfaces were taken. After the first test, the models were re-assembled in a new holder and  5928  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E I  Summary of arc jet sample characteristics, conditions and measured surface temperatures  Calibrated heat flux (W/cm2 )a  Heat flux as referenced  Steady state surface temp.b (  Model #  Density (g/cc)  in this text  Pressure (MPa)     C)  Durationc (s)  FF-61-1  9.59  285  Low condition  0.005  1690  1200  FF-61-2  9.57  350  High condition  0.007  2360  1200  aCold wall heat flux as referenced to a 76 mm dia. Cu hemisphere. b Temperature measured during second exposure. cConsists of two 10 min exposures.  re-instrumented, to enable in-depth or back-face temperature measurements. In this current work we have limited our discussions to the surface temperature measurements only.  3. Results and discussion  3.1. Processing  Table II shows the results of the chemical analysis performed on the raw, milled and hot pressed powders. In the HfB2 powder, the highest impurities were C, O and Zr. Typically, Hf and Zr are found together in natural ores and Hf is a by-product of Zr refining. As a result of their common chemistry and source, both Hf and Zr tend to retain certain impurity levels of each other [14]. Metal borides are typically made by reaction of the purified metal oxide with boronand carbon-containing species. This reaction is capable of yielding large quantities of powder, but it may contain by-products, such as borates. Currently, it is unclear what level of contaminant causes a significant effect on the high temperature properties of UHTCs. As shown in Table II, C and O account for the largest quantity of contaminants in the powders used for this work. In the SiC powder, O is the highest impurity. The milled powder mixture shows higher concentrations of W and O than before. This is due to contamination introduced during the milling process, which used tungsten carbide (WC) milling media. XRD scans of the raw HfB2 powder, the mixed HfB2 /SiC powder and the as-hot-pressed sample (Fig. 6) clearly show that HfB2 peaks dominate. In all of the scans it is difficult to discern SiC peaks because of X-ray absorbance by the HfB2 phase. Present in the scans of the raw and milled powders are small peaks indicating the presence of HfO2 and some ZrO2 . After hot pressing, some HfC was detectable in the solid sample.  T A B L E I I  Elemental characterization of raw powders, milled powders and hot pressed materials  Chemical analysis (wt%)  Sample  B  C  Co  Hf  O  Si  W  Zr  Other  aRaw HfB2 bRaw SiC  10.49  0.11  -  88.7  0.44  -  -  0.29  0.014  -  29.8  -  -  0.49  69.7  -  -  0.048  Milled HfB2 /SiC Hot pressed  9.61  2.02  0.04  82.7  0.89  3.59  0.83  0.37  0.006  9.7  1.81  0.02  83.4  0.03  3.86  0.75  0.40  0.01  Theoretical HfB2 Theoretical SiC  10.8  -  -  89.2  -  -  -  -  -  -  30.0  -  -  -  70.0  -  -  -  Theoretical HfB2 /SiC mix  10.1  2.0  -  83.2  -  4.7  -  -  -  aRaw HfB2 - Al = 0.014%. aRaw SiC - Al = 0.012%, Fe = 0.014%, Ni = 0.011%, V = 0.011%.  The median particle size of the HfB2 and SiC in the as-received powders is 4.1 and 1.6 µm respectively (Table III). After milling for 2 h, the median particle size of the HfB2 -20 vol% SiC mixture is 2.1 µm. It is uncertain whether further reduction in particle size would benefit processing or material properties. A milling time of 2 h was chosen simply to minimize contamination from oxygen and milling media. To further minimize contamination, cyclohexane was used as a milling solvent and an inert atmosphere was used to dry the powders. The resulting particle size of the milled HfB2 /SiC powders is also shown in Table III. Some oxidation of the newly formed surfaces was inevitable during processing, and the O content of the milled powders increased to 0.89 weight percent. As mentioned previously, mixed powders are hot pressed at 2200 C and 25 MPa for 1 h in graphfoillined graphite dies. In this work, both 25 mm diameter ×50 mm tall and 50 mm diameter × 50 mm tall cylindrical billets were hot pressed. These billets required about 250 and 1000 g of powder, respectively. The theoretical density for the HfB2 -20 vol% SiC material is 9.54 g/cm3 assuming a HfB2 density of 11.12 g/cm3 and a SiC density of 3.2 g/cm3 . Fig. 7 shows the densities of thin disks cut along the billet axis as a function of position within the billet, from top to bottom. The disks were 25 or 50 mm in diameter and 1-2 mm thick. Plotted in this figure is bulk density as a function of position, for two 25 mm diameter billets, #52 and #53, and a 50 mm diameter billet, #64. As can be from about 9.5 to 9.75 g/cm3 which is 2% higher seen for the 25 mm diameter billets, the densities range than theoretical. Three possibilities were looked at in trying to explain the variations in density. First, assuming that all of the oxygen contamination of the SiC was in the form SiO2 , calculations indicate that at the measured levels, 0.89 wt%, oxygen in the SiC would     5929  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 6 X-ray scans of raw powders, mixed/milled powders and an as-hot-pressed monolithic billet.  affect final composite density less than half a percent. Next, we assumed the 0.83 wt% tungsten was present as tungsten carbide. With that assumption, calculated density increases less than half a percent, which again is not the observed difference. Density measurements using the Archimedes method indicate that there is no open porosity within the billets. Examination of polished cross-sections using SEM also do not indicate the presence of closed porosity within the billets. Instead, image analysis from several SEM micrographs suggests that the higher than theoretical densities are attributable to a slight loss of SiC during hot pressing. Image analysis clearly shows that a loss of 2-3% SiC occurs, and revised volume fractions of SiC correlate well with observed densities. Additional measurements are required, however, to obtain sufficient statistics. To accommodate the larger diameter and mass of the samples, the hot pressing profile used for the 50 mm diameter billets was modified slightly from that used for the 25 mm diameter billets, by increasing ramp and dwell times. For the single 50 mm billet that was sliced and inspected, Fig. 7 indicates that there is a gradual increase in the density of the larger billet from top to bottom of the sample. This is probably due to sub-optimal processing of the larger billets, because a limited number of these 50 mm billets have been manufactured to date.  T A B L E I I I  Particle size analysis for raw starting powders and pow ders that were attritor milled for 2 h  Particle size (m)  d (0.1)  d (0.5)  d (0.9)  1.2  0.4  0.7  4.1  1.6  2.1  18.3  3.9  4.7  Sample  Raw HfB2 Raw SiC  Milled HfB2 /SiC  5930  Figure 7 Plot of density vs. position within a billet. There remains a  noticeable gradient axially within 50 mm diameter billets that does not  appear to exist in 25 mm billets.  Chemical analysis of the raw HfB2 powder indicates the presence of some C in the raw powder, which may be excess from the carbothermal reduction of HfO2 to yield HfB2 (in the presence of B and C). Residual HfO2 and possible ZrO2 are also present to some degree in the raw powders and even more so after additional oxygen contamination during milling. Chemical analysis of the hot-pressed material shows the continued presence of W remaining from the milling media. Nevertheless, there is a significant drop in O content, from 0.89% in the milled powders to 0.03% in the hot-pressed material. It is acknowledged that the O contents have been  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E  I V Room temperature mechanical properties measured of  the most recent Ames-produced 50 mm billets  Vickers hardness  Modulus  Billet #  61  62  63  65  SHARP-B2  (GPa)  17 ± 0.3 18 ± 0.6 17 ± 0.6 17 ± 0.7 21 ± 1.0  (GPa)  536 ± 5 547 ± 3 546 ± 8 549 ± 8 524 ± 45  Strength  (MPa)  406 ± 57 411a 415 ± 81 453 ± 46 356 ± 91  K IC  √  (MPa  m)  -  - 4.1 ± 0.2 4.2a -  aOnly one sample was  tested,  therefore,  standard deviation was not  calculated.  bers are comparable to the values observed in recent materials manufactured for the SHARP-B2 program (19 to 21 GPa), as shown in Table IV. SEM images of the indents (Fig. 9) show that the crack propagation in the two materials is substantially different. In the materials developed at ARC, the cracks tend to propagate through the grains (intra-granularly) whereas cracks in the previous materials tend to propagate between grains (inter-granularly). The elastic modulus of the Ames materials, 540 GPa, is slightly higher than previous material measurements (Table IV). The higher modulus is likely the result of a slight reduction in the SiC content of these materials compared to the SHARP-B2 materials, due to the loss of SiC during hot-pressing, as discussed previously. The room temperature fracture toughness of these materials averages 4.1 MPa m, which is comparable to the toughness observed in other monolithic ceramics. Average strengths of the Ames materials are greater than 400 MPa, which is again slightly higher than in  √  Figure 8 Representative microstructures of a 50 mm × 50 mm billet showing uniform microstructures (low magnification) and grain struc ture of the SiC phase (black) within the HfB2 phase (gray), (high magnification).  measured only on a small number of raw and milled powders and hot-pressed billets, so the statistical variation in O content from lot to lot and billet to billet is not known. Nonetheless, from this limited data set, it is believed that the drop in O content observed in the hot-pressed billet is a significant result that occurs during each hot pressing run, and does not represent a statistical anomaly. XRD also indicates the presence of a slight amount of HfC in the final hot pressed material. The densities of the flat face arc jet models, machined from a single 50 mm diameter billet, were very similar −9.59 and 9.57 g/cc respectively. Fig. 8 shows SEM micrographs of polished cross sections from hot-pressed HfB2 with a nominal 20% SiC content. The dark phase in the images is SiC and the lighter phase is HfB2 . Fig. 8a is a low magnification image showing the uniform distribution of the SiC phase in the HfB2 matrix. Fig. 8b is a higher magnification image showing the part, SiC grains range from 2-10 µm, while the HfB2 individual SiC particles and HfB2 grains. For the most grains are typically 10-20 µm.  3.2. Properties  Vickers hardness values for the materials processed during this work are between 17 and 18 GPa. These num Figure 9 Hardness  indents  in  Ames  and  SHARP-B2  HfB2 -SiC  composites.  5931  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 10 Plot of strength vs. billet number. Circles indicate data points  where failure was due to graphite inclusions imparted into the materials  during processing. Inset image is of a bend bar fracture surface showing  the observed graphite flaws.  Figure 11 Plot of strength vs. probability of survival ( Ps ) that yields the modulus of statistical reliability for ceramics or Weibull modulus.  previous materials (Table IV). In general, during this study, the average strengths and strength distributions have improved with processing experience, i.e., with the increased numbers of billets processed. This is most clearly seen in Fig. 10, which plots the strengths of individual test specimens as a function of billet number. As this figure shows, the strengths for billet #65 have a very narrow distribution, except for two outlying specimens with considerably lower strengths. These outliers and the lowest valued strength specimen from billet #63 all showed large C defects on the fracture surfaces (see inset image in Fig. 10). We believe that these C defects are graphfoil flakes that were accidentally introduced into the specimens during the die packing process. If these defects are eliminated from consideration, the strengths and strength distributions for billet #65 appear very promising. The effect of these large extrinsic graphite defects on the material’s strength can be seen in the Weibull plot in Fig. 11. The first line plots materials from billets #61, #62, #63 and #65, all 4-pt bend specimens (billet #64 was not mechanically tested as it was sliced to evaluate axial density gradients). The Weibull modulus for this set of specimens is 7. If we take this same data set, but eliminate the bars that failed from the large extrinsic graphfoil defects observed by fractography, the Weibull modulus increases to 11. Going further, if we plot the data from billet #65, without the bend bars that failed from the graphfoil defects, the modulus improves to 15, which is very good for a monolithic ceramic. The coefficient of thermal expansion (CTE) of the current materials, as a function of temperature, is plotted in Fig. 12. For HfB2 -20 vol% SiC, the CTE increases from 3.5 × 10 −6 /K at −100 C to 7 × 10 −6 /K at 1600 C. As expected, the CTE falls between that of pure HfB2 and pure SiC. The UHTCs relatively large CTE, compared to more traditional thermal protection materials such as silica-based tiles, provides a num       5932  Figure 12 Thermal expansion of ARC HfB2 /SiC in comparison to pure HfB2 and pure SiC.  ber of design challenges that must be considered. The large CTE makes the materials less resistant to thermal shock. The large CTE also makes it difficult for designers trying to account for the CTE’s of multi-part components, where the goal is to prevent cracking due to part impingement as a result of expansion during heating. Thermal conductivity measurements were performed on materials from 25 and 50 mm diameter billets, to investigate the consistency from billet to billet. Fig. 13 shows the conductivity as a function of temperature from −100 to 2000 C, for both ARC specimens. This figure shows that there is a slight billet-to-billet variation but the trends are the same: a slight increase in conductivity with temperature. However, there is a significant difference in trends between the current ARC materials and materials from previous eras, such as those from ManLabs and the SHARP-B2 program,     \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 14 Optical  images of flat-face arc jet specimens after their first  and second exposures. Each exposure was for 10 min.        these conditions: it had only a dusty white appearance but no significant amount of oxidation. This is not unexpected at temperatures around 1700 C, which are close to the use temperatures for SiC, employed as a coating on reinforced carbon-carbon (RCC), and the specimen is in a regime where SiO2 is relatively stable. At the high condition, sample 61-2, model temperatures initially followed the profile for sample 61-1 (low) and then increased steadily after 300 s. The surface temperature histories for model FF-61-2 (high condition) for both arc jet exposures are shown in Fig. 16. This figure shows the temperature measurements for both the 1-color and 2-color optical pyrometers. During the initial exposure, the surface temperature rapidly rose and almost leveled off at 1800 C for 300 s. Then the surface temperature began gradually to rise, reaching a maximum temperature between 2200 and 2300 C. During the second exposure, the surface temperature rose rapidly to about the same temperature the model exthe first exposure: 2300 perienced at the end of C. After the first exposure, there was a slight amount of oxidation apparent on the model. During the second exposure, the oxide layer thickened substantially. This observation is consistent, as the surface temperature during the duration of the 2nd run was between 2300 and 2400 C, while it only reached these temperatures at the end of the first exposure. As demonstrated in Fig. 16, the 1-color and the 2-color pyrometers were in close agreement during both exposures of this model. In general, there do not appear to be significant changes in the morphology of the oxide layers radially across the diameter of the model surfaces. A comparison of the images in Fig. 14 shows that the scale of           5933  Figure 13 Plot of thermal conductivity for the Ames materials in com parison to previous work on HfB2 materials [11, 12].  which are also shown in the figure. Materials manufactured for the previous works have significantly higher conductivities (>100 W/mK) at room temperature and the conductivities rapidly decrease with increasing temB2 materials decreases from 130 W/mK at room temperature. For instance, the conductivity in the SHARPperature to 70 W/mk at 1800 C. In contrast, the materials manufactured at ARC have a room temperature conductivity of 40 W/mK and the conductivity increases to 50 W/mK at 1800 C. The reason for the difference in conductivities is not currently understood, but work is continuing in this area in hopes of increasing the conductivity of the Ames materials.        3.3. Arc jet testing        Fig. 14 provides a comparison of the surfaces of the two flat face models after each arc jet exposure. At the low condition and after two exposures, there was little change in the visual appearance of the surface, no measurable weight gain was observed, and steady state surface temperatures were 1690-1750 C. At the high condition, there was some evidence of surface oxidation after the first exposure and there was a corresponding weight loss of 0.5%. After the second exposure, surface oxidation is more evident, a total weight loss of 1% was measured, and steady state surface temperatures in excess of 2300 C were measured. After the second exposure at the high condition, spalling of the surface oxide layers was evident. As confirmed by XRD analysis, the oxide surfaces consist of HfO2 . The measured weight loss is primarily due to loss of BO, B2O2 and SiO because at these high temperatures and low pressures, B2O3 and SiO2 are not stable and are volatilized [15]. Fig. 15 shows the temperature profiles measured by the 1-color pyrometer from each sample during their first tion) the temperature of the surface reached 1750 and second run. For sample 61-1 (low condiC and remained constant. The surface of the model did not change significantly during the two exposures at     \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 15 Summary of 1st and 2nd arc jet exposure temperature histories for both high and low condition flat-face samples.  Figure 16 Summary of 1st and 2nd arc jet exposure temperature histories for sample 61-2, high condition. After 300 s, the temperature can be seen to rise to levels measured during run 2.  the oxide increases with increasing heat flux from low to high condition. On a microscopic scale, the size of what appear as oxide bumps is significantly larger for sample 61-2 than for sample 61-1 (Fig. 17). In previous studies by ManLabs, the oxidation of the HfB2 /SiC samples was observed to leave three distinct regions within the material [3]. The first region comprises the surface oxide, primarily composed of porous HfO2 . Below that is a SiC-depleted zone, where the SiC has been oxidized away (active oxidation), leaving behind a porous HfB2 matrix. Below the depletion zone is the base material. The flat face models tested in this series were crosssectioned to determine if a similar structure had formed. Fig. 18 shows an SEM image of a cross section of sample 61-2, tested at the high condition. The image clearly shows a white oxide layer below which there is a SiC depleted region of porous HfB2 matrix. Similar results  were seen on the arc jet model tested at the low condition, sample 61-1, but the SiC depletion layer was considerably smaller. Table V lists the depths of the oxide layers and the depletion layers observed for each of the flat face models. Oxidation of the SiC clearly plays a significant role in the behavior of these materials during arc jet testing. At 20 volume percent SiC, if the SiC particles  T A B L E V Summary of post-arc jet sample characteristics, including  SiC depletion and oxide layer thicknesses  Model #  qstag  (W/cm2 )  Weight  change (%)  \\x01 Thickness (µm)a  Oxide  layer  (µm)  SiC depletion  layer (µm)  FF-61-1  FF-61-2  285  350  0.0 −1.0  +20 +95  70  340  2  740  aRefers to the overall thickness increase of the sample.  5934  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 17 SEM images of the surfaces of FF-61-1 and 61-2 after two 10-min exposures. Images show the bubbly, porous structure of the surface  oxide.     are assumed to be small spheres randomly distributed throughout the HfB2 matrix, the amount of SiC should be above the percolation threshold. This means that the SiC particles form a network that is interconnected in three dimensions. Polished cross sections do not show this interconnectivity, but this is a result of examining a 3-dimensional network in two dimensions. Additional quantitative stereology, including serial sectioning, is needed to verify SiC interconnectivity. At the low condition, at temperatures approaching 1700 C and at low pressures, the SiC phase forms a passivating layer that helps maintain steady state temperatures. At the high condition, however, the SiC undergoes a transition from passive to active oxidation and the protective SiO2 layer on the SiC is removed as SiO, CO and CO2 . When the gases SiO, BO and B2O2 escape, they pass through the HfO2 , leaving behind channels of porosity. Because HfO2 is not passivating and does not completely seal the surface of a sample, oxygen can diffuse through the porous oxide channels, and newly exposed SiC surfaces are then subject to further oxidation. Cross sections of the models show that the depth of the SiC depletion layer increases with increasing heat flux. It is hypothesized that the extensive formation of the SiC depletion zone is directly related to the 3D interconnected nature of the SiC phase, allowing continuous oxidization of the SiC during the arc jet tests. During testing, the models were probably in a temperature-pressure regime where active oxida tion occurred, drawing more and more SiC out of the sample. The rise in the temperature during the last 300 seconds of the first exposure of sample 61-2 is probably due to a combination of effects. Initially the UHTC surface is oxidized forming a protective SiO2 layer. At the high condition however, the SiO2 layer rapidly becomes unstable and is removed, leading to the formation of a significant HfO2 layer. The formation of HfO2 could lead to increased surface catalycity and reduced surface emittance. Since the conductivity of the oxide is lower than that of the diboride, the surface temperature may also increase as oxide thickness grows. As the test continues, BO, B2O2 and SiO continue to be evolved. However, the data in Table V suggest that further temperature increase is limited by the continued oxidation of the SiC phase, as identified by the increased depth of the SiC depletion zone at the high condition. More arc jet testing is required to improve our understanding of the oxidation mechanisms within these materials in high temperature oxidizing environments. Additional testing is also required to verify the reproducibility of the current results. Extended duration testing is needed over a wider range of conditions to provide more kinetic data for these high temperature oxidation reactions. These tests are needed to develop a model to describe, in detail, oxidation of the HfB2 /SiC system. The effect of stagnation pressure on the oxidation rates also needs to be evaluated.  5935  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 18 Collage of SEM images of a cross section of flat face model  61-2. Note the bubbly porous structure of the oxide above a porous SiC  depleted region. This cross section is the result of two 10-min exposures  at the high condition.  4. Conclusion  Dense HfB2 -20 vol%SiC Ultra High Temperature Ceramics (UHTCs) were produced at NASA Ames by hot pressing. The hardness of the Ames materials was similar to, but slightly less than, that which was measured previously (SHARP-B2). The flexural strengths (at room temperature) were greater than those measured in previous studies. Work is ongoing in the evaluation of the high temperature mechanical properties of these systems. Thermal properties of the Ames materials were compared to previous work (SHARP-B2) and the Ames materials had lower conductivities in comparison to heritage results. Work is continuing to understand the cause of the reduced conductivity in the Ames materials. Arc jet testing of two flat face models provided insight into the oxidation behavior of HfB2 -20 vol% SiC  5936     materials. The models were tested at two different heat fluxes, a high and a low condition, for two 10-minute durations each. It was shown that, during the low condition, the passive oxidation of the SiC portion of these materials plays a role in maintaining steady state surface temperatures at around 1700 C. At the higher heat eled off, at 2400 flux, temperature increased significantly and then levC, after 300 s. Increased temperatures, at the high condition, were attributed to the presence of a thick porous HfO2 layer, which changes the surface catalycity, emittance and conductivity. A further increase in temperature is limited by the continued, active oxidation, of SiC. This is evidenced by the presence of a SiC depleted zone, found when performing post-arc jet analysis of the test specimens.     Acknowledgements  The authors would like to thank the following for their contributions in support of this work: the NASA Ames Research Center Arc Jet Crew and Joe Olejniczak, Ed Martinez, Ricardo Olivares, Tom Squire, and Y.K. Chen of NASA Ames Research Center. This work was performed in collaboration with Lockheed Martin under the NASA Space Launch Initiative program. We would like to also thank Paul Sannes, Jerry Draper and Frank Kustas of Lockheed Martin for their support. Portions of this work were also performed under NASA contract NAS2-99092 to ELORET. We would like to thank Jerry Ridge and Mairead Stackpoole of ELORET for their support.  References  K A U F M A N and E .  1. L . of Boride Compounds for High Temperature Applications,” RTD C L O U G H E R T Y ,  “Investigation  V .  TRD-N69-73497, Part XXXVII, ManLabs Inc., Cambridge, MA,  Dec. 1963.  2.  Idem., “Investigation of Boride Compounds for Very High Temper ature Applications,” RTD-TRD-N63-4096, Part III, ManLabs Inc.,  Cambridge, MA, March 1966. 3. E . “Research and Development of Refractory Oxidaton Resistant Di C L O U G H E R T Y ,  K A L I S H and E .  P E T E R S ,  V .  D .  T .  borides,” AFML-TR-68-190 (1968).  4. A . K . K U R I A K O S E and J .  L . M A R G R A V E , J. Electrochem.  Soc. (1964) 827.  5. 6. L .  J . B . B E R K O W I T Z M A T T U C K , ibid. 113 (1966) 908.  K A U F M A N , B E R K O W I T Z M A T T U C K ,  V .  E .  C L O U G H E R T Y  Trans.  TMS-AIME  and  239  J .  B .  (1967)  458.  7. E . V . C L O U G H E R T Y , R .  L .  P O B E R and L . K A U F M A N ,  ibid. 242 (1968) 1077. 8. M . M .  O P E K A , Z A Y K O S K I and S .  A .  I .  G .  T A L M Y ,  E .  J . W U C H I N A ,  J . C A U S E Y , J. Europ. Ceram. Soc. 19  J .  (1999) 2405.  9.  10.  S .  E .  L E V I N E , S I N G H and J . F . M O N T E V E R D E ,  O P I L A , M .  H A L B I G ,  J .  K I S E R , M .  S A L E M , ibid. 22 (2002) 2757. and S .  B E L L O S I  A .  G U I C C I A R D I ,  ibid. 22 (2002) 279. 11. D . M . A . B A L B O N I , “Arc Heated Facilities,” Advanced Hypersonic Test Facilities , edited by F. Lu and D. Marren (AIAA Progress Series  F E L D E R M A N ,  S H O P E and J .  S M I T H ,  E .  L .  F .  J .  12.  (2002) Vol. 198.  P .  K O L O D Z I E J ,  and D . Flight Demonstration of a Sharp Ultra-High Temperature Ceramic  K E E S E ,  S A L U T E  “First  L .  J .  Nosetip,” NASA TM-112215, Dec. 1997.  \\x0c', '13. R . and H . ity Characterization of Refractory Materials Under High Velocity  P E R K I N S ,  K A U F M A N  N E S O R ,  “Stabil L .  Atmospheric Flight Conditions,” Part  III, Vol  II, ManLabs  Inc.,  Cambridge, MA, 1969.  14.  J .  B .  H E D R I C K , “Zirconium and Hafnium,” U.S. Geological  Survey Minerals Yearbook (1999) p. 86.2.  ULTRA-HIGH TEMPERATURE CERAMICS  15. R .  H .  L A M O R E A U X ,  D .  L .  H I L D E N B R A N D  and L .  B R E W E R , J. Phys. Chem. Ref. Data 16 (1987) 419.  Received 5 November 2003 and accepted 23 April 2004  5937  \\x0c']"
},{
  "_id": 221,
  "PDF": "Processing, properties and arc jet oxidation of hafnium diboride-silicon carbide ultra high temperature ceramics.pdf",
  "Text": "['ULTRA-HIGH TEMPERATURE CERAMICS  J O U R N A L O F M A T E R I A L S S C I E N C E 3 9 (2 0 0 4 ) 5925 - 5937  Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics  M . G A S C H ELORET Corporation, 690 W. Fremont Ave., Sunnyvale, CA 94087, USA  D . E L L E R B Y , E . I R B Y , S . B E C KM A N NASA, Ames Research Center, MS 234-1, Moffett Field, CA 94035, USA  M . G U SM A N ELORET Corporation, 690 W. Fremont Ave., Sunnyvale, CA 94087, USA  S . J O H N S O N NASA, Ames Research Center, MS 234-1, Moffett Field, CA 94035, USA E-mail: Sylvia.M.Johnson@nasa.gov  The processing and properties of HfB2 -20 vol%SiC ultra high temperature ceramics were examined. Dense billets were fabricated by hot-pressing raw powders in a graphite element furnace for 1 h at 2200 C. Specimens were then tested for hardness, mechanical strength, thermal properties and oxidation resistance in a simulated re-entry environment. Thermal conductivity of the current materials was found to be less than previous work had determined while the strength was greater. Oxidation testing of two flat-face models was conducted, at two conditions, for two 10-min durations each. It was concluded that passive oxidation of SiC plays a role in determining the steady-state surface temperatures below 1700 C. Above 1700 C, temperatures are controlled by the properties of a thick HfO2 layer C(cid:2) 2004 Kluwer Academic Publishers and active oxidation of the SiC phase.  1.  Introduction     Ceramic borides, such as hafnium diboride (HfB2 ) and zirconium diboride (ZrB2 ), are members of a family of materials with extremely high melting temperatures which have been referred to as Ultra High Temperature Ceramics (UHTCs). UHTCs constitute a class of promising materials for use in high temperature applications, such as sharp leading edges on future generations of reentry vehicles, because of their high melting points and relatively good oxidation resistance in reentry environments [1]. Because of the extremely high melting temperatures of HfB2 is 3300 of these diboride materials (i.e., melting temperature C as shown in Fig. 1) hot pressing at temperatures >2000 C and at high pressures is required to produce dense materials. The most extensive collection of work conducted on diborides of Hf and Zr was performed by ManLabs in the 1960’s and 70’s under contract with the Air Force. Some of the early work by ManLabs used very high pressure (1-2 GPa) hot pressing to densify the pure diborides [1]. In addition to the pure diborides, ManLabs investigated the influence of a variety of additives, including C and SiC, on the processing, oxidation resistance and thermal shock resistance of Hf and Zr diborides [2]. Their work, in both static furnace experiments and arc jet testing, showed that the addition of SiC improved the     oxidation resistance of Hf and Zr diborides. SiC additions ranging from 5-50 vol% were studied. The work indicated that 20 vol% of a SiC particulate additive resulted in the optimum oxidation resistance [3]. The SiC additions were also found to improve the processing of the materials by reducing the maximum temperature and pressure required to densify the materials and by reducing grain growth of the diboride phase [3]. A number of other studies have been conducted on the oxidation resistance of Hf and Zr diboride materials of various compositions, but these have focused primarily on static or flowing air studies under ambient pressure, and few studies have evaluated their oxidation behavior in simulated reentry environments [4-10]. Static or flowing air oxidation studies under ambient conditions may not provide an accurate representation of a material’s behavior in the unique environment encountered during re-entry. In addition, the relative oxidation resistance of two materials measured in a typical furnace oxidation study may not pertain when testing the same two materials in a reentry environment. Arc jet facilities, on the other hand, provide sustained conditions that are similar to the aerothermal environment experienced during reentry. The testing results are used to understand the thermal performance of materials and systems under controlled aerothermal heating conditions. Arc jet testing is also used to validate  0022-2461  C(cid:2) 2004 Kluwer Academic Publishers  5925  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  properties of the materials have been evaluated and compared to the results from heritage materials of similar compositions (SHARP-B2). Finally, a preliminary investigation of the oxidation behavior of this composition in a simulated reentry environment was begun and testing was conducted at NASA Ames Research Center (ARC) in the Interactive Heating Facility (IHF).  2. Experimental procedure  The raw powders used in this work were −325 mesh HfB2 from Cerac Inc. and 1-2 µm SiC from H.C. Starck Inc. Analysis of the impurities present in the materials was conducted by NSL Analytical Services (Cleveland, OH). Samples were analyzed, using direct current plasma emission spectroscopy to determine metallic impurities, and using inert gas fusion to determine oxygen and carbon impurities. Analysis of the crystalline phases present in the raw and mixed powders was performed with X-ray Diffraction (Scintag X-ray diffractometer) using Cu Kα radiation. Particle size of the raw and milled powders was measured using laser light scattering (Malvern). Light scattering patterns were analyzed with computer software that used the Mie theory. Prior to hot pressing, the HfB2 -20 vol% SiC raw powders were wet-milled with WC milling media in either a planetary mill (Fritsch) or an attrition mill (Union Carbide). The milled powders were carefully dried to prevent phase segregation between the HfB2 (ρ = 11.12 g/cc) and the SiC (ρ = 3.2 g/cc). After drying, the powders were loaded into either 25 or 50 mm diameter graphite dies, lined with graphfoil. Hot pressing cycles were initiated in a graphite element resistance furnace (Thermal Technologies) with vacuum levels of <200 milliTorr. Above 1600 C, the partial pressure of carbon in the furnace increased significantly and degraded the vacuum. Consequently the furnace was back-filled with one atmosphere of inert gas (argon or helium), to preserve the graphite element and insulation. Typical furnace conditions required for densification were 2200 C for 1 h at 25 MPa. The density of the hot pressed billets and test specimens was measured using the Archimedes method and He pycnometry (Micromeritics Accupyc 1330). The phases present in the sintered materials were characterized using XRD. After the samples were crosssectioned and polished to a 1-µm finish, the microstructure of consolidated samples was characterized using optical microscopy and Scanning Electron Microscopy (FEI ESEM 30) with EDX analysis. Mechanical and thermal property test specimens were prepared with diamond tooling, and ground to a final surface finish in accordance with ASTM C1161 (Chand Kare Technical Associates, Worcester, MA). Vickers hardness was determined using indentations created by a Shimadzu HSV-30. Hardness values were calculated from the average length of the diagonal lines across an indentation made using a load of 49 N applied over a 15 s interval. Hardness indents were made in five groups of 3, starting at the outside edge and working towards the center of a sample, to provide a qualitative evaluation of any radial variation in the material microstructure.        Figure 1 Hafnium diboride phase diagram.  the numerical models of materials and systems that are used as design tools [11]. Arc jet testing provides the best ground-based simulation of the reentry environment. Nevertheless, there are a number of differences between the arc jet environment and the reentry environment that must be accounted for when designing an arc jet experiment, and when interpreting the data. For example, surface catalycity can play a more significant role during arc jet testing than it does in flight, because a higher proportion of the air molecules are dissociated in the arc jet environment. NASA Ames began work on UHTCs in the early 1990’s. In 1997 and 2000 Ames demonstrated the use of UHTC sharp leading edges on the Sharp Hypersonic Aero-thermodynamic Research Probe-Ballistic experiments 1 and 2 (SHARP-B1 and SHARP-B2) [12]. The SHARP-B1 vehicle tested a HfB2 -20 vol%SiC nose tip with a 3.5 mm radius. The SHARP-B2 vehicle tested three different leading edge materials, HfB2 -SiC, ZrB2 SiC and ZrB2 -SiC-C. For the sharp wing leading edge applications envisioned for these materials on future reentry vehicles, in addition to oxidation resistance, high thermal conductivity is desirable. In the family of UHTCs, the diborides typically have the highest conductivities. High conductivity improves a material’s thermal shock resistance by reducing temperature gradients and thermal stresses in the material. A higher conductivity also allows more energy to be conducted away from the stagnation point on the wing leading edge. That energy is subsequently reradiated away from lower temperature regions, along the outer surface of the wing leading edge component. This process may enable the vehicle to operate under higher heat flux conditions, improving its performance. The current study has investigated a method of processing HfB2 -20 vol% SiC particulate composites, using conventional hot pressing. Mechanical and thermal  5926  \\x0c', '      quency method, using a GrindoSonic on 3 × 4 × 45 mm Elastic moduli were measured by the resonance frebars. Two methods were employed to measure flexural 1499, using 25.4 mm diameter ×1 mm thick disks) and strength-bi-axial flexure testing (according to ASTM 4-point bend testing (3 × 4 × 45 mm bars according to ASTM C1161), with a crosshead speed of 0.5 mm/min. Estimates of the reliability of the materials were made and represented by a Weibull modulus that was calculated following ASTM C1239. Thermal conductivity of the materials was measured in the range -130 to 2000 C (sample size: 1.27 cm dia. × 0.1 cm thick) using the laser flash technique according to ASTM C201. Coefficient of thermal expansion was measured using a pushrod dilatometer in the temperature range of −130 to 1500 3 × 4 × 10 mm) according to ISO 17562. All thermal C (sample size: property measurements were performed by Netzsch Instruments (Boston, MA). Finally, arc jet testing was performed on machined flat face models, to evaluate the oxidation/ablation behavior of the HfB2 /SiC material. Fig. 2 shows a schematic of the types of arc heaters in use at Ames Research Center. The arc heater produces a hightemperature gas stream by combined radiative, conductive and convective heat transfer from a high-voltage DC electric arc discharge to a gas flowing through a cooled column. The facilities at Ames are capable of input power levels of 20 to 60 MW for up to 30-min durations. The facilities can also test a variety of sample sizes and shapes, such as 80 × 80 cm 2-dimensional articles for panel and seal applications, or 3-dimensional models. The Ames Arc Jet facility can also test samples up to 60 cm in diameter, for aero surface control/response applications. Examples of various HfB2  ULTRA-HIGH TEMPERATURE CERAMICS  UHTC based models that have recently been tested at Ames are shown in Fig. 3. A photo of an as-machined flat face arc-jet model is shown in Fig. 4. The model is 25.4 mm in diameter and the overall height is 8 mm. The notch in the base of the model, shown in the lower left hand picture in Fig. 4, is used to pin the model into the holder. Models were placed in SiC-coated graphite holders (shown in the right hand image of the same figure) which enabled test durations in excess of 10 min; an uncoated graphite holder would have allowed only for a few minutes of testing. Models and holders were then attached to a water-cooled arm (sting), as shown in Fig. 5, that moved the models in and out of the plasma stream. A variety of instrumentation was used to calibrate the arc jet conditions and to measure the thermal response of the materials. The cold wall heat flux values, shown in Table I, were measured using a copper Gardon gage and are referenced to a 76 mm diameter hemisphere. However, the hot wall heat flux values at the model’s surface are different from these calibrations, because of differences in model geometry and differences between the catalycity of the models and the catalycity of the copper Gardon gage. Detailed discussion of the heat flux at the surface of the model is beyond the scope of this work. To differentiate the two conditions, for this text, the first will be designated as “low” and the second will be designated as “high” condition. Two 1-color and one 2-color optical pyrometers were used to make surface temperature measurements during the tests. Based on previous experiments conducted by ManLabs, an emittance of 0.65 was assumed for the 1-color pyrometer [13]. In general there was excellent agreement between the different pyrometers, so  Figure 2 Schematic of the Ames Arc Jet facilities, showing the basic construction of an arc jet. Figure also demonstrates the various nozzle types  available for performing testing on 3-dimensional samples or flat panels [13].  5927  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 3 The family of arc jet models (flat-face, cone and wedge) fabricated and tested at NASA Ames.  Figure 4 Pre-test photo of a flat face arc jet sample, top, side and in a SiC coated graphite holder. Graphite pins were used to hold the sample in place.  To prevent the pins from eroding, the ends were coated with a SiC paste (lighter color spots on the holder).  Figure 5 Photo of a flat-face sample during an arc jet run. Sample and holder are shown attached to a water-cooled sting arm, which serves to move  the sample in and out of the flow.  data from each pyrometer is not shown in every figure. Instrumentation (such as thermocouples or optical pyrometers, for making in-depth or back face temperature measurements) was placed inside a cavity within the sting arm, for protection during the experiments.  For each condition shown in Table I, a model was tested twice for 10-min for a total run time of 20 min. After each exposure, the models were weighed and optical images of their surfaces were taken. After the first test, the models were re-assembled in a new holder and  5928  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E I  Summary of arc jet sample characteristics, conditions and measured surface temperatures  Calibrated heat flux (W/cm2 )a  Heat flux as referenced  Steady state surface temp.b (  Model #  Density (g/cc)  in this text  Pressure (MPa)     C)  Durationc (s)  FF-61-1  9.59  285  Low condition  0.005  1690  1200  FF-61-2  9.57  350  High condition  0.007  2360  1200  aCold wall heat flux as referenced to a 76 mm dia. Cu hemisphere. b Temperature measured during second exposure. cConsists of two 10 min exposures.  re-instrumented, to enable in-depth or back-face temperature measurements. In this current work we have limited our discussions to the surface temperature measurements only.  3. Results and discussion  3.1. Processing  Table II shows the results of the chemical analysis performed on the raw, milled and hot pressed powders. In the HfB2 powder, the highest impurities were C, O and Zr. Typically, Hf and Zr are found together in natural ores and Hf is a by-product of Zr refining. As a result of their common chemistry and source, both Hf and Zr tend to retain certain impurity levels of each other [14]. Metal borides are typically made by reaction of the purified metal oxide with boronand carbon-containing species. This reaction is capable of yielding large quantities of powder, but it may contain by-products, such as borates. Currently, it is unclear what level of contaminant causes a significant effect on the high temperature properties of UHTCs. As shown in Table II, C and O account for the largest quantity of contaminants in the powders used for this work. In the SiC powder, O is the highest impurity. The milled powder mixture shows higher concentrations of W and O than before. This is due to contamination introduced during the milling process, which used tungsten carbide (WC) milling media. XRD scans of the raw HfB2 powder, the mixed HfB2 /SiC powder and the as-hot-pressed sample (Fig. 6) clearly show that HfB2 peaks dominate. In all of the scans it is difficult to discern SiC peaks because of X-ray absorbance by the HfB2 phase. Present in the scans of the raw and milled powders are small peaks indicating the presence of HfO2 and some ZrO2 . After hot pressing, some HfC was detectable in the solid sample.  T A B L E I I  Elemental characterization of raw powders, milled powders and hot pressed materials  Chemical analysis (wt%)  Sample  B  C  Co  Hf  O  Si  W  Zr  Other  aRaw HfB2 bRaw SiC  10.49  0.11  -  88.7  0.44  -  -  0.29  0.014  -  29.8  -  -  0.49  69.7  -  -  0.048  Milled HfB2 /SiC Hot pressed  9.61  2.02  0.04  82.7  0.89  3.59  0.83  0.37  0.006  9.7  1.81  0.02  83.4  0.03  3.86  0.75  0.40  0.01  Theoretical HfB2 Theoretical SiC  10.8  -  -  89.2  -  -  -  -  -  -  30.0  -  -  -  70.0  -  -  -  Theoretical HfB2 /SiC mix  10.1  2.0  -  83.2  -  4.7  -  -  -  aRaw HfB2 - Al = 0.014%. aRaw SiC - Al = 0.012%, Fe = 0.014%, Ni = 0.011%, V = 0.011%.  The median particle size of the HfB2 and SiC in the as-received powders is 4.1 and 1.6 µm respectively (Table III). After milling for 2 h, the median particle size of the HfB2 -20 vol% SiC mixture is 2.1 µm. It is uncertain whether further reduction in particle size would benefit processing or material properties. A milling time of 2 h was chosen simply to minimize contamination from oxygen and milling media. To further minimize contamination, cyclohexane was used as a milling solvent and an inert atmosphere was used to dry the powders. The resulting particle size of the milled HfB2 /SiC powders is also shown in Table III. Some oxidation of the newly formed surfaces was inevitable during processing, and the O content of the milled powders increased to 0.89 weight percent. As mentioned previously, mixed powders are hot pressed at 2200 C and 25 MPa for 1 h in graphfoillined graphite dies. In this work, both 25 mm diameter ×50 mm tall and 50 mm diameter × 50 mm tall cylindrical billets were hot pressed. These billets required about 250 and 1000 g of powder, respectively. The theoretical density for the HfB2 -20 vol% SiC material is 9.54 g/cm3 assuming a HfB2 density of 11.12 g/cm3 and a SiC density of 3.2 g/cm3 . Fig. 7 shows the densities of thin disks cut along the billet axis as a function of position within the billet, from top to bottom. The disks were 25 or 50 mm in diameter and 1-2 mm thick. Plotted in this figure is bulk density as a function of position, for two 25 mm diameter billets, #52 and #53, and a 50 mm diameter billet, #64. As can be from about 9.5 to 9.75 g/cm3 which is 2% higher seen for the 25 mm diameter billets, the densities range than theoretical. Three possibilities were looked at in trying to explain the variations in density. First, assuming that all of the oxygen contamination of the SiC was in the form SiO2 , calculations indicate that at the measured levels, 0.89 wt%, oxygen in the SiC would     5929  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 6 X-ray scans of raw powders, mixed/milled powders and an as-hot-pressed monolithic billet.  affect final composite density less than half a percent. Next, we assumed the 0.83 wt% tungsten was present as tungsten carbide. With that assumption, calculated density increases less than half a percent, which again is not the observed difference. Density measurements using the Archimedes method indicate that there is no open porosity within the billets. Examination of polished cross-sections using SEM also do not indicate the presence of closed porosity within the billets. Instead, image analysis from several SEM micrographs suggests that the higher than theoretical densities are attributable to a slight loss of SiC during hot pressing. Image analysis clearly shows that a loss of 2-3% SiC occurs, and revised volume fractions of SiC correlate well with observed densities. Additional measurements are required, however, to obtain sufficient statistics. To accommodate the larger diameter and mass of the samples, the hot pressing profile used for the 50 mm diameter billets was modified slightly from that used for the 25 mm diameter billets, by increasing ramp and dwell times. For the single 50 mm billet that was sliced and inspected, Fig. 7 indicates that there is a gradual increase in the density of the larger billet from top to bottom of the sample. This is probably due to sub-optimal processing of the larger billets, because a limited number of these 50 mm billets have been manufactured to date.  T A B L E I I I  Particle size analysis for raw starting powders and pow ders that were attritor milled for 2 h  Particle size (m)  d (0.1)  d (0.5)  d (0.9)  1.2  0.4  0.7  4.1  1.6  2.1  18.3  3.9  4.7  Sample  Raw HfB2 Raw SiC  Milled HfB2 /SiC  5930  Figure 7 Plot of density vs. position within a billet. There remains a  noticeable gradient axially within 50 mm diameter billets that does not  appear to exist in 25 mm billets.  Chemical analysis of the raw HfB2 powder indicates the presence of some C in the raw powder, which may be excess from the carbothermal reduction of HfO2 to yield HfB2 (in the presence of B and C). Residual HfO2 and possible ZrO2 are also present to some degree in the raw powders and even more so after additional oxygen contamination during milling. Chemical analysis of the hot-pressed material shows the continued presence of W remaining from the milling media. Nevertheless, there is a significant drop in O content, from 0.89% in the milled powders to 0.03% in the hot-pressed material. It is acknowledged that the O contents have been  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  T A B L E  I V Room temperature mechanical properties measured of  the most recent Ames-produced 50 mm billets  Vickers hardness  Modulus  Billet #  61  62  63  65  SHARP-B2  (GPa)  17 ± 0.3 18 ± 0.6 17 ± 0.6 17 ± 0.7 21 ± 1.0  (GPa)  536 ± 5 547 ± 3 546 ± 8 549 ± 8 524 ± 45  Strength  (MPa)  406 ± 57 411a 415 ± 81 453 ± 46 356 ± 91  K IC  √  (MPa  m)  -  - 4.1 ± 0.2 4.2a -  aOnly one sample was  tested,  therefore,  standard deviation was not  calculated.  bers are comparable to the values observed in recent materials manufactured for the SHARP-B2 program (19 to 21 GPa), as shown in Table IV. SEM images of the indents (Fig. 9) show that the crack propagation in the two materials is substantially different. In the materials developed at ARC, the cracks tend to propagate through the grains (intra-granularly) whereas cracks in the previous materials tend to propagate between grains (inter-granularly). The elastic modulus of the Ames materials, 540 GPa, is slightly higher than previous material measurements (Table IV). The higher modulus is likely the result of a slight reduction in the SiC content of these materials compared to the SHARP-B2 materials, due to the loss of SiC during hot-pressing, as discussed previously. The room temperature fracture toughness of these materials averages 4.1 MPa m, which is comparable to the toughness observed in other monolithic ceramics. Average strengths of the Ames materials are greater than 400 MPa, which is again slightly higher than in  √  Figure 8 Representative microstructures of a 50 mm × 50 mm billet showing uniform microstructures (low magnification) and grain struc ture of the SiC phase (black) within the HfB2 phase (gray), (high magnification).  measured only on a small number of raw and milled powders and hot-pressed billets, so the statistical variation in O content from lot to lot and billet to billet is not known. Nonetheless, from this limited data set, it is believed that the drop in O content observed in the hot-pressed billet is a significant result that occurs during each hot pressing run, and does not represent a statistical anomaly. XRD also indicates the presence of a slight amount of HfC in the final hot pressed material. The densities of the flat face arc jet models, machined from a single 50 mm diameter billet, were very similar −9.59 and 9.57 g/cc respectively. Fig. 8 shows SEM micrographs of polished cross sections from hot-pressed HfB2 with a nominal 20% SiC content. The dark phase in the images is SiC and the lighter phase is HfB2 . Fig. 8a is a low magnification image showing the uniform distribution of the SiC phase in the HfB2 matrix. Fig. 8b is a higher magnification image showing the part, SiC grains range from 2-10 µm, while the HfB2 individual SiC particles and HfB2 grains. For the most grains are typically 10-20 µm.  3.2. Properties  Vickers hardness values for the materials processed during this work are between 17 and 18 GPa. These num Figure 9 Hardness  indents  in  Ames  and  SHARP-B2  HfB2 -SiC  composites.  5931  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 10 Plot of strength vs. billet number. Circles indicate data points  where failure was due to graphite inclusions imparted into the materials  during processing. Inset image is of a bend bar fracture surface showing  the observed graphite flaws.  Figure 11 Plot of strength vs. probability of survival ( Ps ) that yields the modulus of statistical reliability for ceramics or Weibull modulus.  previous materials (Table IV). In general, during this study, the average strengths and strength distributions have improved with processing experience, i.e., with the increased numbers of billets processed. This is most clearly seen in Fig. 10, which plots the strengths of individual test specimens as a function of billet number. As this figure shows, the strengths for billet #65 have a very narrow distribution, except for two outlying specimens with considerably lower strengths. These outliers and the lowest valued strength specimen from billet #63 all showed large C defects on the fracture surfaces (see inset image in Fig. 10). We believe that these C defects are graphfoil flakes that were accidentally introduced into the specimens during the die packing process. If these defects are eliminated from consideration, the strengths and strength distributions for billet #65 appear very promising. The effect of these large extrinsic graphite defects on the material’s strength can be seen in the Weibull plot in Fig. 11. The first line plots materials from billets #61, #62, #63 and #65, all 4-pt bend specimens (billet #64 was not mechanically tested as it was sliced to evaluate axial density gradients). The Weibull modulus for this set of specimens is 7. If we take this same data set, but eliminate the bars that failed from the large extrinsic graphfoil defects observed by fractography, the Weibull modulus increases to 11. Going further, if we plot the data from billet #65, without the bend bars that failed from the graphfoil defects, the modulus improves to 15, which is very good for a monolithic ceramic. The coefficient of thermal expansion (CTE) of the current materials, as a function of temperature, is plotted in Fig. 12. For HfB2 -20 vol% SiC, the CTE increases from 3.5 × 10 −6 /K at −100 C to 7 × 10 −6 /K at 1600 C. As expected, the CTE falls between that of pure HfB2 and pure SiC. The UHTCs relatively large CTE, compared to more traditional thermal protection materials such as silica-based tiles, provides a num       5932  Figure 12 Thermal expansion of ARC HfB2 /SiC in comparison to pure HfB2 and pure SiC.  ber of design challenges that must be considered. The large CTE makes the materials less resistant to thermal shock. The large CTE also makes it difficult for designers trying to account for the CTE’s of multi-part components, where the goal is to prevent cracking due to part impingement as a result of expansion during heating. Thermal conductivity measurements were performed on materials from 25 and 50 mm diameter billets, to investigate the consistency from billet to billet. Fig. 13 shows the conductivity as a function of temperature from −100 to 2000 C, for both ARC specimens. This figure shows that there is a slight billet-to-billet variation but the trends are the same: a slight increase in conductivity with temperature. However, there is a significant difference in trends between the current ARC materials and materials from previous eras, such as those from ManLabs and the SHARP-B2 program,     \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 14 Optical  images of flat-face arc jet specimens after their first  and second exposures. Each exposure was for 10 min.        these conditions: it had only a dusty white appearance but no significant amount of oxidation. This is not unexpected at temperatures around 1700 C, which are close to the use temperatures for SiC, employed as a coating on reinforced carbon-carbon (RCC), and the specimen is in a regime where SiO2 is relatively stable. At the high condition, sample 61-2, model temperatures initially followed the profile for sample 61-1 (low) and then increased steadily after 300 s. The surface temperature histories for model FF-61-2 (high condition) for both arc jet exposures are shown in Fig. 16. This figure shows the temperature measurements for both the 1-color and 2-color optical pyrometers. During the initial exposure, the surface temperature rapidly rose and almost leveled off at 1800 C for 300 s. Then the surface temperature began gradually to rise, reaching a maximum temperature between 2200 and 2300 C. During the second exposure, the surface temperature rose rapidly to about the same temperature the model exthe first exposure: 2300 perienced at the end of C. After the first exposure, there was a slight amount of oxidation apparent on the model. During the second exposure, the oxide layer thickened substantially. This observation is consistent, as the surface temperature during the duration of the 2nd run was between 2300 and 2400 C, while it only reached these temperatures at the end of the first exposure. As demonstrated in Fig. 16, the 1-color and the 2-color pyrometers were in close agreement during both exposures of this model. In general, there do not appear to be significant changes in the morphology of the oxide layers radially across the diameter of the model surfaces. A comparison of the images in Fig. 14 shows that the scale of           5933  Figure 13 Plot of thermal conductivity for the Ames materials in com parison to previous work on HfB2 materials [11, 12].  which are also shown in the figure. Materials manufactured for the previous works have significantly higher conductivities (>100 W/mK) at room temperature and the conductivities rapidly decrease with increasing temB2 materials decreases from 130 W/mK at room temperature. For instance, the conductivity in the SHARPperature to 70 W/mk at 1800 C. In contrast, the materials manufactured at ARC have a room temperature conductivity of 40 W/mK and the conductivity increases to 50 W/mK at 1800 C. The reason for the difference in conductivities is not currently understood, but work is continuing in this area in hopes of increasing the conductivity of the Ames materials.        3.3. Arc jet testing        Fig. 14 provides a comparison of the surfaces of the two flat face models after each arc jet exposure. At the low condition and after two exposures, there was little change in the visual appearance of the surface, no measurable weight gain was observed, and steady state surface temperatures were 1690-1750 C. At the high condition, there was some evidence of surface oxidation after the first exposure and there was a corresponding weight loss of 0.5%. After the second exposure, surface oxidation is more evident, a total weight loss of 1% was measured, and steady state surface temperatures in excess of 2300 C were measured. After the second exposure at the high condition, spalling of the surface oxide layers was evident. As confirmed by XRD analysis, the oxide surfaces consist of HfO2 . The measured weight loss is primarily due to loss of BO, B2O2 and SiO because at these high temperatures and low pressures, B2O3 and SiO2 are not stable and are volatilized [15]. Fig. 15 shows the temperature profiles measured by the 1-color pyrometer from each sample during their first tion) the temperature of the surface reached 1750 and second run. For sample 61-1 (low condiC and remained constant. The surface of the model did not change significantly during the two exposures at     \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 15 Summary of 1st and 2nd arc jet exposure temperature histories for both high and low condition flat-face samples.  Figure 16 Summary of 1st and 2nd arc jet exposure temperature histories for sample 61-2, high condition. After 300 s, the temperature can be seen to rise to levels measured during run 2.  the oxide increases with increasing heat flux from low to high condition. On a microscopic scale, the size of what appear as oxide bumps is significantly larger for sample 61-2 than for sample 61-1 (Fig. 17). In previous studies by ManLabs, the oxidation of the HfB2 /SiC samples was observed to leave three distinct regions within the material [3]. The first region comprises the surface oxide, primarily composed of porous HfO2 . Below that is a SiC-depleted zone, where the SiC has been oxidized away (active oxidation), leaving behind a porous HfB2 matrix. Below the depletion zone is the base material. The flat face models tested in this series were crosssectioned to determine if a similar structure had formed. Fig. 18 shows an SEM image of a cross section of sample 61-2, tested at the high condition. The image clearly shows a white oxide layer below which there is a SiC depleted region of porous HfB2 matrix. Similar results  were seen on the arc jet model tested at the low condition, sample 61-1, but the SiC depletion layer was considerably smaller. Table V lists the depths of the oxide layers and the depletion layers observed for each of the flat face models. Oxidation of the SiC clearly plays a significant role in the behavior of these materials during arc jet testing. At 20 volume percent SiC, if the SiC particles  T A B L E V Summary of post-arc jet sample characteristics, including  SiC depletion and oxide layer thicknesses  Model #  qstag  (W/cm2 )  Weight  change (%)  \\x01 Thickness (µm)a  Oxide  layer  (µm)  SiC depletion  layer (µm)  FF-61-1  FF-61-2  285  350  0.0 −1.0  +20 +95  70  340  2  740  aRefers to the overall thickness increase of the sample.  5934  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 17 SEM images of the surfaces of FF-61-1 and 61-2 after two 10-min exposures. Images show the bubbly, porous structure of the surface  oxide.     are assumed to be small spheres randomly distributed throughout the HfB2 matrix, the amount of SiC should be above the percolation threshold. This means that the SiC particles form a network that is interconnected in three dimensions. Polished cross sections do not show this interconnectivity, but this is a result of examining a 3-dimensional network in two dimensions. Additional quantitative stereology, including serial sectioning, is needed to verify SiC interconnectivity. At the low condition, at temperatures approaching 1700 C and at low pressures, the SiC phase forms a passivating layer that helps maintain steady state temperatures. At the high condition, however, the SiC undergoes a transition from passive to active oxidation and the protective SiO2 layer on the SiC is removed as SiO, CO and CO2 . When the gases SiO, BO and B2O2 escape, they pass through the HfO2 , leaving behind channels of porosity. Because HfO2 is not passivating and does not completely seal the surface of a sample, oxygen can diffuse through the porous oxide channels, and newly exposed SiC surfaces are then subject to further oxidation. Cross sections of the models show that the depth of the SiC depletion layer increases with increasing heat flux. It is hypothesized that the extensive formation of the SiC depletion zone is directly related to the 3D interconnected nature of the SiC phase, allowing continuous oxidization of the SiC during the arc jet tests. During testing, the models were probably in a temperature-pressure regime where active oxida tion occurred, drawing more and more SiC out of the sample. The rise in the temperature during the last 300 seconds of the first exposure of sample 61-2 is probably due to a combination of effects. Initially the UHTC surface is oxidized forming a protective SiO2 layer. At the high condition however, the SiO2 layer rapidly becomes unstable and is removed, leading to the formation of a significant HfO2 layer. The formation of HfO2 could lead to increased surface catalycity and reduced surface emittance. Since the conductivity of the oxide is lower than that of the diboride, the surface temperature may also increase as oxide thickness grows. As the test continues, BO, B2O2 and SiO continue to be evolved. However, the data in Table V suggest that further temperature increase is limited by the continued oxidation of the SiC phase, as identified by the increased depth of the SiC depletion zone at the high condition. More arc jet testing is required to improve our understanding of the oxidation mechanisms within these materials in high temperature oxidizing environments. Additional testing is also required to verify the reproducibility of the current results. Extended duration testing is needed over a wider range of conditions to provide more kinetic data for these high temperature oxidation reactions. These tests are needed to develop a model to describe, in detail, oxidation of the HfB2 /SiC system. The effect of stagnation pressure on the oxidation rates also needs to be evaluated.  5935  \\x0c', 'ULTRA-HIGH TEMPERATURE CERAMICS  Figure 18 Collage of SEM images of a cross section of flat face model  61-2. Note the bubbly porous structure of the oxide above a porous SiC  depleted region. This cross section is the result of two 10-min exposures  at the high condition.  4. Conclusion  Dense HfB2 -20 vol%SiC Ultra High Temperature Ceramics (UHTCs) were produced at NASA Ames by hot pressing. The hardness of the Ames materials was similar to, but slightly less than, that which was measured previously (SHARP-B2). The flexural strengths (at room temperature) were greater than those measured in previous studies. Work is ongoing in the evaluation of the high temperature mechanical properties of these systems. Thermal properties of the Ames materials were compared to previous work (SHARP-B2) and the Ames materials had lower conductivities in comparison to heritage results. Work is continuing to understand the cause of the reduced conductivity in the Ames materials. Arc jet testing of two flat face models provided insight into the oxidation behavior of HfB2 -20 vol% SiC  5936     materials. The models were tested at two different heat fluxes, a high and a low condition, for two 10-minute durations each. It was shown that, during the low condition, the passive oxidation of the SiC portion of these materials plays a role in maintaining steady state surface temperatures at around 1700 C. At the higher heat eled off, at 2400 flux, temperature increased significantly and then levC, after 300 s. Increased temperatures, at the high condition, were attributed to the presence of a thick porous HfO2 layer, which changes the surface catalycity, emittance and conductivity. A further increase in temperature is limited by the continued, active oxidation, of SiC. This is evidenced by the presence of a SiC depleted zone, found when performing post-arc jet analysis of the test specimens.     Acknowledgements  The authors would like to thank the following for their contributions in support of this work: the NASA Ames Research Center Arc Jet Crew and Joe Olejniczak, Ed Martinez, Ricardo Olivares, Tom Squire, and Y.K. Chen of NASA Ames Research Center. This work was performed in collaboration with Lockheed Martin under the NASA Space Launch Initiative program. We would like to also thank Paul Sannes, Jerry Draper and Frank Kustas of Lockheed Martin for their support. Portions of this work were also performed under NASA contract NAS2-99092 to ELORET. We would like to thank Jerry Ridge and Mairead Stackpoole of ELORET for their support.  References  K A U F M A N and E .  1. L . of Boride Compounds for High Temperature Applications,” RTD C L O U G H E R T Y ,  “Investigation  V .  TRD-N69-73497, Part XXXVII, ManLabs Inc., Cambridge, MA,  Dec. 1963.  2.  Idem., “Investigation of Boride Compounds for Very High Temper ature Applications,” RTD-TRD-N63-4096, Part III, ManLabs Inc.,  Cambridge, MA, March 1966. 3. E . “Research and Development of Refractory Oxidaton Resistant Di C L O U G H E R T Y ,  K A L I S H and E .  P E T E R S ,  V .  D .  T .  borides,” AFML-TR-68-190 (1968).  4. A . K . K U R I A K O S E and J .  L . M A R G R A V E , J. Electrochem.  Soc. (1964) 827.  5. 6. L .  J . B . B E R K O W I T Z M A T T U C K , ibid. 113 (1966) 908.  K A U F M A N , B E R K O W I T Z M A T T U C K ,  V .  E .  C L O U G H E R T Y  Trans.  TMS-AIME  and  239  J .  B .  (1967)  458.  7. E . V . C L O U G H E R T Y , R .  L .  P O B E R and L . K A U F M A N ,  ibid. 242 (1968) 1077. 8. M . M .  O P E K A , Z A Y K O S K I and S .  A .  I .  G .  T A L M Y ,  E .  J . W U C H I N A ,  J . C A U S E Y , J. Europ. Ceram. Soc. 19  J .  (1999) 2405.  9.  10.  S .  E .  L E V I N E , S I N G H and J . F . M O N T E V E R D E ,  O P I L A , M .  H A L B I G ,  J .  K I S E R , M .  S A L E M , ibid. 22 (2002) 2757. and S .  B E L L O S I  A .  G U I C C I A R D I ,  ibid. 22 (2002) 279. 11. D . M . A . B A L B O N I , “Arc Heated Facilities,” Advanced Hypersonic Test Facilities , edited by F. Lu and D. Marren (AIAA Progress Series  F E L D E R M A N ,  S H O P E and J .  S M I T H ,  E .  L .  F .  J .  12.  (2002) Vol. 198.  P .  K O L O D Z I E J ,  and D . Flight Demonstration of a Sharp Ultra-High Temperature Ceramic  K E E S E ,  S A L U T E  “First  L .  J .  Nosetip,” NASA TM-112215, Dec. 1997.  \\x0c', '13. R . and H . ity Characterization of Refractory Materials Under High Velocity  P E R K I N S ,  K A U F M A N  N E S O R ,  “Stabil L .  Atmospheric Flight Conditions,” Part  III, Vol  II, ManLabs  Inc.,  Cambridge, MA, 1969.  14.  J .  B .  H E D R I C K , “Zirconium and Hafnium,” U.S. Geological  Survey Minerals Yearbook (1999) p. 86.2.  ULTRA-HIGH TEMPERATURE CERAMICS  15. R .  H .  L A M O R E A U X ,  D .  L .  H I L D E N B R A N D  and L .  B R E W E R , J. Phys. Chem. Ref. Data 16 (1987) 419.  Received 5 November 2003 and accepted 23 April 2004  5937  \\x0c']"
},{
  "_id": 222,
  "PDF": "Processing, sintering and oxidation behavior of SiC fibers reinforced ZrB2 composites.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 32 (2012) 1933-1940  Processing, sintering and oxidation behavior of SiC fibers reinforced ZrB2 composites  Diletta Sciti  , Laura Silvestroni  ∗  CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy  Available online 25 November 2011  Abstract  Borides and carbides of early transition metals are considered a class of promising materials for several applications, the most appealing ones being in the aerospace and energy sectors. The present work is mostly focused on toughening of UHTCs through incorporation of SiC chopped fibers. Mechanical properties of reinforced composites are compared to those of un-reinforced, whiskerand particle-reinforced materials and the effect of different kinds of sintering aids is studied. Addition of fibers allows toughness to be increased from 3-4 MPa m1/2 (for un-reinforced materials) to 5.0-6.2 MPa m1/2 . The high temperature behavior is also investigated both in air furnace and in arc jet facility. Eventually, a paragraph is dedicated to potential of UHTCs as sunlight absorbers for future solar concentrating systems operating in the high temperature regime. © 2011 Elsevier Ltd. All rights reserved.  Keywords: Borides; Composites; Fibres; Toughness and toughening; Corrosion  1.   Introduction  Zirconium diboride-based materials are currently considered a class of promising materials for several applications,  in particular  in  the aerospace sector.  In  the  last five years,  research has focused on  the fabrication of dense composites possessing high strength  (500-1000 MPa).1-11 However,  the  low  fracture toughness remains one of the major concerns for the application of these materials under severe environmental conditions. From data reported in the literature it is evident that often the addition of particles does not represent an effective strategy for a major toughness improvement. In case of SiC particles, for example, it has been shown that residual tensile stresses are developed in the ZrB2 matrix, due  to  the difference of  thermal expansion coefficient between SiC and ZrB2 .7 As a result, particle-reinforced ZrB2-SiC materials are often as brittle as other ZrB2 -based composites even if they display significant increase of hardness and strength. Spherical reinforcement can be efficaciously substituted by elongated reinforcement. The potential advantages of elongated secondary phases over particulate-reinforced  systems  include  ∗  Corresponding author.  E-mail address: diletta.sciti@istec.cnr.it (D. Sciti).  0955-2219/$ - see front matter © 2011 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2011.10.032     more effective toughening mechanisms such as enhanced crack deflection and load-carrying capability. In addition, other toughening mechanisms,  such  as  crack bridging  and pullout,  are possible. Indeed, significant increases of fracture toughness were obtained through addition of SiC whiskers,12,13 carbon fibers14 and more recently SiC chopped fibers.15-17 Oxidation of refractory diborides  in air at elevated  temperatures has  limited  their applications. One of  the commonest additives used to improve the oxidation protection of ZrB2 is silicon carbide18-20 as it forms a coherent SiO2 -based glass layer on the surface of the ceramic. This continuous scale provides passive oxidation protection in the intermediate temperature regime (1100-1600 C). The aim of this work is to evaluate the efficacy of SiC chopped fibers  as potential  reinforcement  for ZrB2 analysing  several aspects, such as: the impact of the sintering technique, the effect of  the  sintering aid,  the high  temperature  stability,  including both conventional tests in air furnace up to 1700 C and arc jet tests  to simulate re-entry conditions at  temperatures exceeding 2000 C. Furthermore, the performance of SiC fiber-reinforced composites is compared to that of whiskeror particle-reinforced composites at  room, high  temperature and  in oxidising environment. Indeed, whether or not SiC chopped fibers can be as effective as SiC particles during oxidation is still an unexplored issue. A short paragraph at the end of the paper is dedicated to potential application of UHTCs in the solar energy sector.              \\x0c', '1934   D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940  2. Experimental  2.1. Preparation of SiC chopped fibers reinforced materials  \\u242em   Commercial raw materials were used to prepare the ceramic composites: ZrB2 Grade B  (H.C. Starck, Germany), particle size range 0.1-8  \\u242em; SiC HI Nicalon chopped fibers, with composition (wt%) Si:C:O = 62:37:0.5. The dimensions are: 14  \\u242em diameter, 1 mm length. As sintering additives the following pow␣-Si3N4 Baysinid (Bayer, ders were used in amount 5-10 vol%:  Germany); ZrSi2 -F (Japan New Metals Co., LTD, Osaka, Japan); MoSi2 <2  (Aldrich, Steinbeim, Germany). The SiC fibers content varied from 0 to 30 vol%. Particular care needs to be paid for processing of fiber-reinforced materials. Homogenization of fibers and powders was accomplished by conventional ball milling for 24 using ZrO2 media in absolute ethanol. However, compaction was generally more difficult  than for conventional powder mixtures. The slurries were dried  in a rotary evaporator and the powder mixtures underwent a slow debonding cycle in a graphite  furnace at 500 C  for 60 min  in flowing Argon. The powder-fiber mixtures were sintered  in a graphite mould using a hot pressing machine at 1650, 1700 and 1750-1900 C for compositions doped with ZrSi2 , Si3N4 and MoSi2 , respectively, with applied pressure of 40-50 MPa and holding  time of 10 min. The microstructures were analyzed using scanning electron microscopy (SEM, Cambridge S360, Cambridge, UK) and energy dispersive spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK) on fractured and polished surfaces. For comparison, SiC whiskerand particle-reinforced ZrB2 composites were produced. Details on  their preparation are given elsewhere.11,15        2.2. Mechanical characterization  ×  ×  Vickers microhardness  (HV1.0) was measured by  indentation with a 9.8 N  load  in a hardness  tester Zwick model 3212. Fracture  toughness (KIc ) was evaluated using chevron notched beams (CNB) in flexure. The test bars, 25 mm   2 mm   2.5 mm (length by width by thickness, respectively), were notched with a  0.1 mm-thick  diamond  saw;  the  chevron-notch  tip  depth and average side  length were about 0.12 and 0.80 of  the bar thickness,  respectively. The specimens were  fractured using a semi-articulated silicon carbide four-point fixture with a  lower span of 20 mm and an upper span of 10 mm on a universal testing machine Instron 1195. The specimens,  three for each composite, were  loaded with a crosshead speed of 0.05 mm/min. The “slice model” equation of Munz et al.21 was used  to calculate KIc . On the same machine and with the same fixture, the flexural strength (σ ) was measured at room temperature and at 1200 C in air.     2.3. High temperature characterization  The oxidation  tests were carried out  in a bottom-up  loading furnace at 1200, 1500 and 1700 C  for 30 min  in static air on various compositions containing either SiC chopped fibers or whiskers or particles. Specimens were  located  in  the  furnace        jet   ×  ×  Arc   when the maximum temperature was reached and then removed and  air  quenched  after  the  exposure  time. Mass  and  bars’ dimensions were measured  before  and  after  the  oxidation. Microstructural modifications  induced  by  oxidation  were analyzed by SEM-EDS. Flexural  strength  tests  in  four point fixture were conducted on 25 mm   2 mm   2.5 mm bars after oxidation at 1700 C. tests: A composite containing 20 vol% fibers was selected for the realization of a sharp leading edge with a radius of curvature <0.5 mm. The experiments were conducted  in an arc-jet plasma wind tunnel at the Department of Aerospace Engineering - University of Naples “Federico II”, and involved high enthalpy supersonic flows  in air with a nominal Mach number (M) around 3. The sharp wedge was  locked upon an alumina holder at a distance of 10 mm from the nozzle exit. The surface temperature of the wedge was continuously measured by a twocolour pyrometer (Infratherm ISQ5, Impac Electronic GmbH, Germany)  focused onto a position as close as possible  to  the edge of  the sample. At  the same  time,  the distribution of  the surface  temperature over  the wedge was monitored using an infrared thermo-camera (FLIR Thermacam P40, spectral range 8-14  \\u242em). Microstructural modifications  induced by oxidation were then analyzed by SEM-EDS.  3. Results and discussion  3.1. Microstructure and properties of materials containing chopped fibers  For composites containing up to 20 vol% of fibers, few or no residual porosity was generally observed  in  the microstructure and the fiber dispersion into all the matrices was homogeneous, Fig. 1a. A significant fraction of residual porosity (8%) was noticed  in  the matrix containing 30 vol% of fibers  (Fig. 1b), either  in  the boride matrix or at  the matrix/fiber  interface.  It is  apparent  that  increasing  the  reinforcement  content makes increasingly difficult  to properly process  the powder mixtures to achieve a dense matrix and a good fiber dispersion. In all the composites a notable  reduction of  the fibers dimension  (maximum  length around 300  \\u242em) was observed, due  to both  the action of milling media and the applied pressure during sintering. Secondary phases were observed in the sintered microstructures, including residual ZrO2 particles, isolated pockets of BN, Zr-Si phases, SiO2 and a borosilicate glass containing Zr-Si-B-N-O distributed along grain boundaries.15,16 As for the fiber-matrix interface (Fig. 1c)  it can be observed  that a very complex system of reaction phases developed: around the original SiC fiber core, a Si-C-O rim was formed probably due  to migration of intergranular phase  towards  the fiber surface. The  latter often reacted with  the  liquid phase deriving  from  the sintering aids and/or gathered various  impurities present  in  the matrix, especially ZrO2 . Fig. 1d shows  the crack propagation  features of these composites: the crack very often crosses the fibers without appreciable deviation due to the strong matrix/fiber interface. The  effect of SiC  chopped fibers on mechanical properties of ZrB2 composites  is  shown  in Fig. 2  for  the  system containing Si3N4 as  sintering agent. Fracture  toughness was  \\x0c', 'D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940   1935  Fig. 1. ZrB2 + SiCf + Si3N4 composites showing (a)  30 vol% fiber, showing porosity in the boride matrix, (c) fiber/matrix interface, and (d) crack path obtained by a 9.8 N Vickers indentation.  typical microstructure of samples containing 20 vol% fiber, (b)   typical microstructure of samples containing  improved significantly by  the addition of fibers, up  to 20 vol% (5.3 and 5.6 MPa m0.5 ) in comparison with the reference materials  (3.7 MPa m0.5 )  (Fig. 2a). The  toughness of a  reinforced composite can be  interpreted as  the  sum of  the unreinforced matrix toughness plus the increment due to toughening mechanisms exerted by the fibers15,16 :  K reinforced Ic  =  Kmatrix Ic  +  \\x01Kfiber  (1)  As crack propagation behavior displayed no pullout, deflection or bridging (Fig. 1c),  it was concluded  that  the dominant toughening mechanism was crack bowing.16 Beside this, a second  toughening  contribution deriving  from  thermal  residual stresses was also considered, due  to mismatch between  thermal expansion coefficients and elastic constants of matrix and fibers.22 As  αmatrix , thermal residual stresses are expected toughness.22 In preto give a negative contribution  to fracture  vious works16 it was shown that the toughening increment due to crack bowing23 was much larger than the negative one from residual stresses  that  is almost negligible,22 resulting  in a net increase of toughness. Also, it was shown that this model is well suited for fiber content up to 20 vol% but fails for a higher content (30 vol%). Indeed, the residual porosity negatively affected  αSiC <   Ic  \\u242em   σ  ) and the reinforcing  both the matrix original toughness (Kmatrix action of the fibers (\\x01Kfiber ). The  flexural  strength  of  fiber-containing  composites decreased  from 600  to 400 MPa  for 10-20 vol% fibers  and 200 MPa  for 30 vol% fibers  (Fig. 2b). The  incorporation of 200-300  long elements  in a fine matrix (ZrB2 mean grain size: 3  \\u242em) very  likely changed  the defect population. Using the Griffith equation,   = 1.3 KIc /C0.5 , the size of the mean critical defect (C) was estimated  to be about 100-200  \\u242em, which suggests  that fibers  themselves could  represent critical flaws. The flexural strength was also tested at 1200 C in air (Fig. 2b). It can be observed  that  the composites nearly maintained  the same RT value (400 MPa). Furthermore, the composites were strengthened in comparison with the un-reinforced material that exhibited a value of 240 MPa at this temperature.     3.2. Effect of the sintering technique  Spark plasma  sintering  (SPS)  is widely  recognized  as  a very effective  sintering  technique  that allows materials  to be densified at  lower  temperatures and  in a shorter  time  than  in conventional  sintering. Recent contributions have  shown  that this  technique permits  the  successful densification of UHTC  Fig. 2. Effect of the fiber content on (a) fracture toughness and (b) 4-pt flexural strength at room temperature and 1200     C in air.      \\x0c', '1936   D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940  Table 1  Comparison of mechanical properties between compositions containing 20 vol% fiber and 5 vol% Si3N4 , obtained by hot pressing (HP) and spark plasma sintering (SPS). KIc : fracture toughness by CNB,  σ RT : 4-pt flexural strength at room temperature,  σ 1200 : 4-pt flexural strength at 1200 C.     Composition,  vol%  Technique   T/t/P,     C/min/MPa  Exp. density, g/cm3  KIc ,  MPa m1/2 5.65 ±  0.30  5.53   0.37   σ RT ,  MPa  σ 1200 ,  MPa  Ref.  ZrB2 + 20 SiC fiber  ZrB2 + 20 SiC fiber   HP  1700/10/50   5.2   410   ± ±   20   340   ± ±   50   [15]  SPS  1500/5/50   4.7   ±  370    20   447    23   [16]  ceramics with reduced content of sintering aids24,25 or with addition of phases (e.g. whiskers or fibers) that are very sensitive to temperatures.13 As already mendegradation at high sintering  tioned, composites processed  in  ISTEC  labs were sintered by hot pressing at 1700 C. For comparison some SPS  tests were conducted  in Arrhenius Lab, Stockholm University, Sweden. The powder mixtures were sintered  in an SPS furnace,  in vacuum  (5-10 Pa). The  sintering  temperature was  in  the  range 1500-1700 C and applied pressure was 50 MPa.16 Microstructural analyses confirmed that composites containing 10-20 vol% fibers were  completely dense  at  temperatures  comparatively lower  than for hot pressing,  i.e. 1500 C. However, despite  the lower sintering temperature and lower holding time, microstructural analysis of  the fiber-matrix  interface confirmed a  large extent of chemical interaction between SiC fibers, intergranular phases and ZrB2 . Mechanical properties of an SPS composite containing 20 vol% fiber are compared  to  those of  the HP composite with  the same composition  in Table 1. The values of fracture toughness of the hot pressed composites were close to  those of  the SPS composites. This  indicates  that, even  if a more efficient  thermal  treatment as SPS was used,  the nature of the matrix/reinforcement interface did not change notably. In particular, the analysis of the interaction between an advancing crack and  the microstructure confirmed  that  the  intergranular wetting phases or  interface reactions prevented any possibility of reinforcement pullout,  independently of  the sintering route. Also  the strength values did not  indicate substantial advantage of the SPS technique over the hot pressing route, at least for the composition tested.16           3.3. Effect of the shape of reinforcement (chopped fibers vs. whiskers and particles)  For  the  sake  of  comparison, SiC whisker and  particlereinforced composites were produced by hot pressing at 1700 C using the same sintering aid, Si3N4 .11,15 The mechanical properties of  the  three kind of composites containing 20 vol% of reinforcement are summarized  in Table 2. Similar  to  the case     of fiber  additions,  fracture  toughness was  improved by  the addition of 20 vol% SiC whiskers  (5.3 MPa m0.5 ). However, incorporation of whiskers did not alter significantly the strength (610 MPa).15 The particle-reinforced composite had a  lower toughness  (4.6 MPa m0.5 )  and higher  strength  (730 MPa),  as expected. Concerning strength test at 1200 C, it can be observed that for particle-reinforced composites,  the strength decreased to 250 MPa, probably due  to deleterious effects associated  to softening of  intergranular phases. In case of fiber and whisker additions, strength values were 340-350 MPa despite  the presence of the same intergranular soft phases. The high-temperature strength  is affected by several  factors: a - softening of amorphous phases that causes a strength decrease; b - change of the defect population induced by oxidation, (the materials were kept about 30 min at a temperature >1000 C) that causes a strength decrease; c - sealing of  small porosities by  the borosilicate glass, that could induce a strength increase. At 1200 C, all the composites underwent effects a-c. The  fact  that whiskerand fiber-containing materials  resulted  in higher strength suggests that these reinforcing elements still have a strengthening effect. This, in turn, could be attributed to the fact that differently from particles, fibers and whiskers underwent a slower or incomplete oxidation and continued to exert a structural function.           3.4. Effect of sintering aid  The choice of  sintering aid affects not only  the  sintering behavior of  the matrix, but also  the extent of reaction between matrix and fibers. A part from Si3N4 , two silicides, namely ZrSi2 and MoSi2 , were tested. The principle behind selection of these silicides is related to their efficacy in promoting the densification and in improving the oxidation resistance of boride matrices.26 ZrSi2 is a relatively low melting point phase (1600 C) and thus densification was completed at 1600-1650 C. For MoSi2 , that is more  refractory, densification  tests were carried out  in  the range 1750-1900 C.17 In all  the cases dense composites were obtained. The analysis of the microstructure revealed that at similar densification temperatures, the addition of MoSi2 promoted higher fiber degeneration  in comparison with Si3N4 or ZrSi2 .           Table 2  Effect of the reinforcement type on mechanical properties of reinforced composites, KIc : fracture toughness by CNB,  σ 1200 : 4-pt flexural strength at 1200 C,  σ ret : retained strength after oxidation at 1700 C/30 min.  σ RT : 4-pt flexural strength at room temperature,        Composition,  vol%  Reinforcement type  Additive,  vol%  KIc ,  MPa m1/2  σ RT ,  MPa  σ 1200 ,  MPa  Ref.   σ ret ,  MPa  ZrB2 + 20 SiC  ZrB2 + 20 SiC  ZrB2 + 20 SiC   Chopped fibers  Si3N4 , 5%  Si3N4 , 5%  Si3N4 , 5%   5.65   ± ± ±   0.30   410   ± ± ±   20   340   ± ± ±   50   [15]   164   ± ± ±   8  Whiskers   5.30    0.30   614    75   352    22   [15]   180    11  Particles   4.55    0.10   730    100   250    10   [11]   220    22  \\x0c', 'D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940   1937  Table 3  Effect of sintering aids on the mechanical properties. KIc : Fracture toughness by CNB,   σ RT : 4-pt flexural strength at room temperature.  Composition,  vol%  ZrB2 ZrB2 + 20 SiC fiber  ZrB2 ZrB2 + 20 SiC fiber  ZrB2 ZrB2 + 20 SiC fiber  ZrB2 + 20 SiC fiber   Additive  Sintering  temperature,     C  Si3N4 Si3N4 ZrSi2 ZrSi2 MoSi2 MoSi2 MoSi2  1700  1700   1600   1600   1750   1750   1900   KIc ,  MPa m0.5  σ RT ,  MPa  3.7 ±  0.1  5.7   0.3   ± ± ± ± ± ±   0.1    0.4    0.6    0.1    0.3   4.3   6.2   3.5   4.8   3.7   600 ± 410   808   385   780   380   -  ± ± ± ± ±   90   20   31   13   87   20     Microstructural features are mirrored in the results of mechanical  properties,  reported  in Table  3. Among  the  different systems, the maximum fracture toughness increase was for the composition containing ZrSi2 as sintering additive, achieving 6.2 MPa m0.5 . No significant  toughness  increase was observed for MoSi2 doped materials:  the fibers degeneration occurring during sintering at 1900 C rendered  the reinforcements completely ineffective. A modest increase of toughness was instead found after sintering treatment at 1750 C. As the toughness of the composites can be interpreted as the sum of the un-reinforced matrix  toughness plus  the  increment due  to  toughening mechanisms exerted by  the fibers (Eq. (1)),  it can be stated  that  the higher  toughness of  the system containing ZrSi2 is due  to  the higher toughness of its respective matrix. In the case of MoSi2 , the  lowest values are both due  to  the  lower  toughness of  the matrix and the fiber degeneration. Finally, one could observe that the change of sintering aid does not have a significant impact on the strength of the composites, being all values around 400 MPa.     3.5. Oxidation behavior        The oxidation  tests were carried out  in a bottom-up  loading air furnace at 1200, 1500 and 1700 C for 30 min on composites containing 20 vol% fibers. After oxidation at 1200 C (not shown), a discontinuous cracked SiO2 scale with ZrO2 particles and damaged fibers was observed on  the surface. After oxidation at 1500 and 1700 C,  islands of ZrO2 in rounded particle or dendrite-like  shape, were  immersed  in a SiO2 continuous scale.27 The polished  sections of fiber-reinforced composites oxidised  at  1500  and  1700 C  are  reported  in Fig.  3a  and b and  they generally  followed  the  typical aspect of oxidised ZrB2-SiC: an external SiO2 layer, an intermediate coarse ZrO2           partially filled with silica and a SiC-depleted  layer above  the unreacted bulk. No SiC-depleted  layer  is present  in  the materials  after  exposure  at 1200 C. After oxidation  at 1500 C carbon residuals were often observed around the fibers (Fig. 3c), which were hardly observed in oxidised SiC particle-reinforced ceramics.28,29 This indicates that active oxidation of SiC fibers proceeds first through the reaction: 2SiC + O2 =   2SiO (g)    2C   +  (2)     instead of the more frequently reported reaction28,29 :  O2 =   SiO (g)    CO (g)   SiC   +  +  (3)  It is interesting to study the effect of SiC reinforcement type on the oxidation behavior. When using chopped fibers, we introduce SiC elements with typical dimensions of 100-300  \\u242em and diameter of 15  \\u242em. The whiskers are typically 30  \\u242em in length \\u242em. In composite containing particles, and diameter less than 1  SiC particles tend to form agglomerates at triple points that do not overcome dimensions around 2-3  \\u242em.  In  terms of weight gain, parallel oxidation experiments demonstrated  that all  the samples displayed similar values, in the range 1200-1700 C.27 However, when a depletion region is formed, in the case of fine SiC particles, a layer with small and fine porosity will form. In the case of fiber additions, the SiC depletion region will contain large cavities. An  intermediate situation  is represented by SiC whiskers. The strength degradation after oxidation at 1700 C due  to  the surface modified  layer  is displayed  in Table 2. As can be  seen,  the absolute value of  retained  strength  is  lower for the composite containing fibers (160 MPa) than for whiskers (180 MPa) and particles  (220 MPa), confirming  that morphology of the SiC phase can have a significant impact on the high temperature performance.        Fig. 3. Polished cross section of ZrB2 + 20SiC fibers + Si3N4 after oxidation at (a) 1500 and (b) 1700     C. (c) Initial stage of fiber oxidation.  \\x0c', '1938   D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940  Fig. 4. External surface of the model (a) before and (b) after the arc-jet test with magnified areas of the modified microstructure.  3.6. Extreme oxidation behavior: arc-jet tests  The design of high-performance hypersonic vehicles requires pointed  fuselage  nose-cones  and wing  leading  edges with very  sharp  profile,  i.e.  radius  of  curvature  in  the  order  of few millimeters. This helps  the vehicle  to  enhance  its performances and maneuvrability by  reducing  the drag and also to  improve  the  crew  safety due  to  an  expanded down  and cross-range capability.30,31 However  for  sharp  leading edges the convective heating  to  the  surface, and hence  the  surface temperature, dramatically  increases as  the  radius of curvature decreases.31 The composition containing 20 vol% fibers and silicon nitride was selected  for  the  realization of a sharp  leading edge with a  radius of curvature  less  than 0.5 mm  (Fig. 4a)  to be  tested in supersonic  regime. The pyrometer was  focused on a 3 mm diameter spot from  the  tip and  the maximum average  temperature reached  in  this region was 1700 C. However, according to computational modeling, on  the surface  tip  the  temperature could reach values as high as 2300-2400 C. The  total  time of exposition  including heating ramps, was 17 min. Fig. 4a and b shows  the appearance of  the model before and after  the  test: the shape  is maintained and  the colour changed from white on the tip (presence of ZrO2 phase) to black at the base (presence of silica-based glass). The most damaged area was obviously        \\u242em   the summit of the sample where large cavities left by SiC fibers evaporation were well visible  together with cracks  (Fig. 4b). The presence of solely ZrO2 phase was observed down to about 300  \\u242em. Large pools of silica containing small ZrO2 precipitates were visible at 400  from  the  surface and at 3 mm the predominant phase was zirconia and a discontinuous  layer of silica. At 6 mm  from  the  tip fibers were still  recognizable, though cracked and damaged, among ZrO2 particles. The cross section of the model was severely damaged by the polishing procedure (Fig. 5a). However the three typical regions observed for ZrB2-SiC composites were recognizable: an outermost cracked zirconia  layer  (40  \\u242em), a ZrB2 -SiO2 interlayer  (140  \\u242em) and a SiC-depleted ZrB2 region (400  \\u242em thick). The appearance of SiC fibers was very similar to the morphology observed for other ZrB2-SiC fibers composites oxidised in air furnace: the core was SiC which was oxygen enriched moving outside, around the fiber a continuous graphite layer was observed (see Fig. 3c). Although the effect of  type of SiC reinforcement needs a more  in-depth investigation, preliminary arc jet tests conducted on a composite containing 20 vol% SiC particles with  the same geometry and in similar conditions did evidence a marked difference  in  the scale morphology  (see Fig. 5b).  In particular  the scale shape is  still  sharp  in  the fiber-reinforced material and flattened  in the particle-reinforced material, suggesting a different extent of ablation between  the  two composites. This behavior can be a  \\x0c', 'D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940   1939  Fig. 5. Cross section of (a) SiC chopped fibers reinforced and (b) SiC particles reinforced ZrB2 .  consequence of a different oxidation rate of fibers and particles and will be object of future investigation.  4. UHTC for solar energy application        Beside  the well known  characteristics  that make UHTCs attractive as thermal protection systems, there is a strong interest in  their applications as sunlight absorbers for solar concentrating systems  that can operate  in  the high  temperature  regime. Electricity generation from solar energy by thermo-mechanical conversion  is currently  limited  in worldwide  implementation, due to the relatively low conversion efficiency and the resulting high cost of the produced electricity. This level of performance and cost is achieved today in solar thermal power plant technologies  that are based on steam cycles at moderate  temperatures (400-550 C). A  real breakthrough may occur  if  the conversion efficiency from sunlight  to electricity can be significantly increased. This would require heat input at temperatures in the range of 1000-1400 C, which  in  turn  faces significant challenges, especially concerning  the key component  that  is  the radiation absorber located at the focus of the concentrator field. The ideal absorber material should possess high emissivity in the solar spectrum (up  to about 2  \\u242em) and  low emissivity outside this range to produce spectral selectivity, reasonable resistance to stresses, high thermal conductivity to smooth local hot spots due to uneven distribution of incident radiation and stability to chemical degradation caused by oxidation over  long operation time at elevated  temperature. Former materials used  for volumetric absorbers  include Al2O3 , SiO2 and SiC. Concerning UHTCs,  there  is some evidence  that  these composites possess intrinsic spectral selectivity,32,33 howa pronounced degree of  ever, the studies of spectral emissivity characteristics of UHTCs performed  so  far  refer  to films and coatings and not  to bulk materials, either porous or dense.  For  this  reason  a  long-term  investigation  about UHTC potentialities as  innovative absorber  for new generation hightemperature concentrating solar power plants, CSP, has recently started  in collaboration with  the  Italian  Institute of National studies34 Optics  (INO-CNR). Preliminary  have  shown  that borides and carbides behave very differently from conventional ceramics such as SiC. Even with some differences among them, all UHTCs display a step-like increase of reflectance from visible to infrared, with presence of a cut-off reflectance wavelength. These characteristics, in combination with other key properties as high thermal and electrical conductivities, superior oxidation resistance up  to 1600 C make unreinforced UHTC monoliths and composites particularly  interesting  for solar applications. Further studies will be devoted to investigation of high temperature optical behavior up to 1200 C or higher.        5. Conclusions and final remarks     The addition of SiC chopped fibers has demonstrated  to be a viable method for improving the fracture toughness of brittle ZrB2 ceramics as values of 5.7-6.2 MPa m1/2 can be reached for compositions containing 20 vol% fibers. As far as oxidation  is concerned the stability of fiber-reinforced composites is comparable to that of SiC particles reinforced composites, either in air furnaces up to 1700 C or in arc jet tests from 1700 to 2400 C, at Mach 3. One drawback for  these composites  is  the decrease of flexural strength which passes from typical 700-800 MPa for un-reinforced or particle  reinforced materials  to 400 MPa  for compositions with 10-20 vol% of fibers. However, it should be mentioned  that  for strength  tests at high  temperature  (at  least up  to 1200 C), flexural strength  is presently higher  than  that of several developed ZrB2 -based materials.  It was shown  that the choice of  the sintering aid  is a key  issue, as  it affects not only  the matrix  sintering  temperature, but also  the  reactivity        \\x0c', '1940   D. Sciti, L. Silvestroni / Journal of the European Ceramic Society 32 (2012) 1933-1940  between matrix and fiber.  In  turn,  the presence of secondary phases deriving from sintering aid can severely affect  the high temperature mechanical and  thermal performance. Among  the possible options, silicon nitride was found to offer the best compromise. Future directions for  increasing UHTCs performance include:  minimization of sintering aid amount for improving high temperature performance up to 1500 C; minimization of RT  strength deterioration  fiber amount and material processing); improvement of thermal shock resistance.     (optimization of   Finally, un-reinforced UHTCs can find a novel field of application as solar absorbers, due to their intrinsic spectral selectivity that  is possessed only by a  limited number of materials. These characteristics,  in  combination with other key properties  as high  thermal and electrical conductivities,  superior oxidation resistance make UHTC monoliths and composites particularly attractive for high temperature solar applications.  Acknowledgements  Authors wish  to  acknowledge R. Savino  for  arc  jet  tests (University of Naples “Federico II”, Napoli, Italy); M. Nygren for SPS  tests  (Arrhenius  lab, Stockolm University, Stockolm, Sweden); E. Sani and L. Mercatelli for optical tests (Institute of National Optics, INO-CNR, Firenze, Italy); S. Guicciardi and C. Melandri for mechanical test and discussion; D. Dalle Fabbriche for the hot pressing and oxidation routes.  References  1. Gasch MJ, Ellerby DT, Johnson SM. Ultra high temperature ceramic com posites. In: Bansal NP, editor. Handbook of ceramic composites. Boston:  Kluwer Academic Publishers; 2005. p. 197-224.  2. Fahrenholtz WG, Hilmas GE, Talmy IG, Zaykoski JA. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90(5):1347-64.  3. Telle R, Sigl LS, Takagi K. Boride-based hard materials. In: Riedel R, editor.  Handbook of ceramic hard materials. Weinheim, Germany: Wiley-VCH;  2000. p. 802-945.  4. Cutler RA. Engineering properties of borides. In: Schneider Jr SJ, editor.  Ceramics and glasses: engineered materials handbook, vol. 4. Materials  Park, OH: ASM International; 1991. p. 787-803.  5. Levine SR, Opila EJ, Halbig MC, Kiser JD, Singh M, Salem JA. Evaluation  of ultra-high temperature ceramics for aeropropulsion use. J Eur Ceram Soc 2002;22:14-5, 2757-67.  6. Chamberlain AL, Fahrenholtz WG, Hilmas GE. High-strength zirconium diboride-based ceramics. J Am Ceram Soc 2004;87(6):1170-2.  7. Watts J, Hilmas GE, Farenholtz WG, Brown D, Clausen B. Stress mea surements in ZrB2 -SiC composites using Raman spectroscopy and neutron diffraction. J Eur Ceram Soc 2010;30:2165-71.  8. Sciti   D,   Guicciardi   S,   Bellosi   A,   Pezzotti   G.   Properties   of   a  pressureless-sintered ZrB2 -MoSi2 ceramic composite. J Am Ceram Soc 2006;89:2320-2.  9. Balbo A, Sciti D. Spark plasma sintering and hot pressing of ZrB2 -MoSi2 ultra-high-temperature ceramics. Mater Sci Eng A 2008;475:108-12.  10. Sciti D, Silvestroni L, Celotti G, Melandri C, Guicciardi S. Sintering and  mechanical properties of ZrB2 -TaSi2 and HfB2 -TaSi2 ceramic composites. J Am Ceram Soc 2008;91:3285-91.  11. Monteverde F, Guicciardi S, Bellosi A. Advanced   in microstructure and  mechanical properties of zirconium diboride based ceramics. Mater Sci Eng A 2003;346:310-9.  12. Zhang P, Hu P, Zhang X, Han   J, Meng S. Processing and characteri zation of ZrB2 -SiCw ultra-high  2008;472:358-62.  temperature ceramics. J Alloys Compd  13. Wang H, Wang CA, Yao X, Fang D. Processing and mechanical properties  of zirconium diboride-based ceramics prepared by spark plasma sintering. J Am Ceram Soc 2007;90:1992-7.  14. Yang F, Zhang X, Han J, Du S. Characterization of hot-pressed short carbon  fiber  reinforced ZrB2 -SiC ultra-high  Alloys Compd 2009;472:395-9.  temperature ceramic composites. J  15. Silvestroni L, Sciti D, Guicciardi S, Melandri C. Toughened ZrB2 -based ceramics with addition of SiC whiskers or chopped fibres. J Eur Ceram Soc 2010;30:2155-64.  16. Guicciardi S, Silvestroni L, Nygren M, Sciti D. Microstructure and tough ening mechanisms  in  spark plasma  sintered ZrB2 ceramics  reinforced by SiC whiskers or SiC chopped fibers. J Am Ceram Soc 2010;93(8):  2384-91.  17. Sciti D, Guicciardi S, Silvestroni L. SiC chopped fibers reinforced ZrB2 : effect of the sintering aid. Scr Mater 2011;64:769-72.  18. Shaffer PTB. An oxidation resistant boride composition. Am Ceram Bull 1962;41(2):96-9.  19. Tripp WC, Davis HH, Graham HC. Effect of an SiC addition on  dation of ZrB2 . Am Ceram Bull 1973;52(8):612-6. 20. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during the  the oxi oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27(6):2495-501.     C. J Eur  21. Munz DG, Shannon Jr JL, Bubsey RT. Fracture toughness calculations from  maximum load in four point bend tests of chevron notch specimens. Int J Fracture 1980;16:R137-41.  22. Taya M, Hayashi S, Kobayashi AS, Yoon HS. Toughening of a particulate reinforced ceramic-matrix composite by  Ceram Soc 1990;73:1382-91.  thermal   residual   stress. J Am  23. Rouxel T, Laurent Y. Fracture characteristics of SiC particle   reinforced  oxynitride glass using chevron-notch  Fracture 1998;91:83-101.  three-point bend   specimens.   Int J  24. Sciti D, Silvestroni L, Nygren M. Spark plasma sintering of Zrand Hf borides with decreasing amounts of MoSi2 as sintering aid. J Eur Ceram Soc 2008;28:1287-96.  25. Sciti D, Guicciardi S, Nygren M. Densification and mechanical behavior  of HfC and HfB2 fabricated by spark plasma sintering. J Am Ceram Soc 2008;91:1433-40.  26. Sciti D, Brach M, Bellosi A. Long-term oxidation behavior and mechanical  strength degradation of a pressureless sintered ZrB2 -MoSi2 ceramic. Scr Mater 2005;53:1297-302.  27. Silvestroni L, Sciti D. Oxidation of ZrB2 ceramics containing SiC as particles, whiskers, or short fibers. J Am Ceram Soc 2011;94(9):2796-9.  28. Fahrenholtz WG.   Thermodynamic   analysis   of   ZrB2 -SiC   oxidation: 2007;90(1):  formation   of   a   SiC-depleted   region.   J   Am Ceram   Soc   143-8.  29. Monteverde F, Bellosi A. Oxidation of ZrB2 -based ceramics  Electrochem Soc 2003;150(11):B552-9.  in dry air. J  30. Squire T, Marschall   J. Material property   requirements   for analysis and  design of UHTC components in hypersonic applications. J Eur Ceram Soc 2010;30:2239-51.  31. Monti R, Savino R, De Stefano Fumo M. Matching flight conditions on  sharp leading edges in plasma wind tunnels. J Thermophys Heat Transfer 2007;21(3):660-4.  32. Bogaerts WF, Lampert CM. Review, Materials  energy conversion. J Mater Sci 1983;18:2847-75.  for photothermal   solar  33. Kennedy   CECE.   Review   of   mid-to   high-temperature   solar   selec tive   absorbers   materials.   National   Renewable   Energy   Laboratory;  2002.  34. Sani E, Mercatelli L, Sansoni P, Silvestroni L, Sciti D. Ultra-high temper ature ceramic absorbers for high-temperature thermodynamic solar plants.  J Renew Sustain Energy, submitted for publication.  \\x0c']"
},{
  "_id": 223,
  "PDF": "Production and Oxidation Resistance of HfB2–30 vol _ SiC composite powders modified with Y3Al5O12.pdf",
  "Text": "['ISSN 0036-0236, Russian Journal of Inorganic Chemistry, 2020, Vol. 65, No. 9, pp. 1416-1423. © Pleiades Publishing, Ltd., 2020. Russian Text © The Author(s), 2020, published in Zhurnal Neorganicheskoi Khimii, 2020, Vol. 65, No. 9, pp. 1274-1282.  INORGANIC MATERIALS  AND NANOMATERIALS  Production and Oxidation Resistance of HfB2-30 vol % SiC  Composite Powders Modif ied with Y3Al5O12  E. P. Simonenkoa, *, N. P. Simonenkoa, I. A . Nagornova, b, V. G. Sevastyanova, and N. T. Kuznetsova  aKurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences, Moscow, 119991 Russia bMendeleev University of Chemical Technology of Russia, Moscow, 125047 Russia *e-mail: ep_simonenko@mail.ru  Received March 16, 2020; revised April 13, 2020; accepted April 27, 2020  Abstract—The HfB2-30 vol % SiC composite powders modif ied with a small (2 and 5 vol %) amount of Y3Al5O12 were synthesized by the sol-gel method. Such a modif ication in producing an ultra-high-temperature ceramic of the corresponding composition optimized the process temperature and increased the oxidation resistance of the ceramic by stabilizing the formed hafnium dioxide in the cubic or tetragonal form. It was determined that the introduced X-ray amorphous phase Y3Al5O12 was distributed as a thin f ilm on particles of hafnium diboride and silicon carbide and led to agglomeration of the latter; no formation of individual Y3Al5O12 particles was observed. Investigation of the thermal behavior of the obtained composite powders in an air f low in the temperature range 20-1200°C showed that the introduction of 5 vol % Y3Al5O12 virtually completely neutralized the effect of the initial surface oxidation of HfB2 in the temperature range 700-800°C, shifted the temperature of the beginning of oxidation and the maximum of the corresponding thermal event from 756 to 808°C, and decreased the intensity of this thermal event.  Keywords: UHTC, composite powder, HfB2, SiC, YAG, sol-gel method, oxidation resistance DOI: 10.1134/S003602362009020X  INTRODUCTION  Numerous  research  groups develop MB2-SiC (M = Zr, Hf) materials proposed to be used during aerodynamic heating to temperatures about ~2500°C [1-15]. Along with a suff icient resistance to oxidation, including oxidation by atomic oxygen [16-20], such ceramic composites have reasonably good mechanical performance, high thermal conductivity (also at elevated temperatures), optical properties suitable for potential applications, and low catalytic activity in exothermic reactions of surface recombination of oxygen and nitrogen atoms, etc. One of  the key properties of ZrB2(HfB2)-SiC ultra-high-temperature ceramics (UHTC) is the ability to endure quite a long exposure to an oxygen-containing atmosphere at elevated temperatures (>1200- 2500°C) owing to the presence of a protective layer of borosilicate glass  formed by  the oxidation of  the ceramic. Depending on  the characteristics of  the material (composition, porosity) and the exposure conditions, this airproof glassy layer can occur both on the surface of the ceramic (at not too high temperatures of the latter: <1700-1850°C), and in the bulk of the porous layer based on the hafnium oxide, a product of the oxidation of HfB2 (at high surface temperature >1900-2000°C). Therefore, investigation of the oxidation resistance of UHTC of a specif ic composi tion (with taking into account various strengthening additives) and the microstructure under certain exposure conditions are extremely important [21-38]. One of the obvious methods to increase the oxidation resistance is to develop methods to obtain maximally dense materials. The structures of zirconium and hafnium diborides and silicon carbide are such that  the production of ceramics based on  them requires quite high temperatures, which leads not only to high energy consumption, but also to impairment of the mechanical performance because of coarsening of grains (f irst of all, of metal diboride) by high-temperature consolidation. Introduction of sintering additives signif icantly reduces the required process temperature. Liquid-phase sintering of oxygen-free ceramics using additives in the system Al2O3-Y2O3, including an additive corresponding to yttrium aluminum garnet Y3Al5O12 (YAG), has long been used to ensure excellent densif ication at not too high temperatures both in pressureless sintering [39-44], and in high-temperature pressing: hot pressing [45-49] or spark plasma sintering [50-54]. The amount, particle size, and the uniformity of the distribution of a sintering additive affect signif icantly not only the densif ication process, but also  the properties of  the obtained material, including the oxidation resistance. However, it was  1416  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1417  noted that a high YAG content in long-term treatment at elevated temperatures and low pressure can lead to degradation of the material [48-56] because of the interaction of the components of the ceramic. Therefore, a most important problem is to minimize the necessary amount of an introduced sintering additive.  One  more  problem  of  the  oxidation  of ZrB2(HfB2)-SiC UHTC by a high-enthalpy air f low is the separation of the oxidized layer because of phase transformations during cooling or reheating of such oxidation products as high-melting ZrO2/HfO2, the thermodynamically stable form of which is the monoclinic crystal lattice. A solution can be modif ication of the ceramic with agents capable of stabilizing hafnium oxide in the cubic or tetragonal system, e.g., with high-melting and low-volatile yttrium oxide [57-59]. For example, it was shown [58] that, in the HfB2- 20 vol % SiC-3 vol % Y2O3 composite (in comparison with the unmodif ied material), a protective surface layer of borosilicate glass is stabilized.  However, a technologically better stabilizing additive could be less high-melting yttrium aluminum garnet. It was demonstrated [60] that, at a temperature >2200°C, from Y3Al5O12, more volatile Al2O3 can be distilled off, and the remaining Y2O3 can be consumed for the stabilization of the product of the oxidation of HfB2. At a lower temperature, the unevaporated alumina  can modify  the  silicate  glass,  specif ically, increase its softening temperatures and the viscosity in the molten state, thus decreasing the oxygen diffusion into the bulk of the material ZrB2(HfB2)-SiC. The eff iciency of the introduction of Y3Al5O12 for improving the oxygen resistance of UHTC has recently been conf irmed [46]: it was concluded that the doping of the ZrB2-SiC ceramics with Al/Y oxides considerably improved their oxidation resistance at the temperature of 1700°C because of a decrease in the activity of silicon dioxide. The ZrB2-SiC-B4C-Y3Al5O12 coatings [61] had a good thermal-shock resistance in cyclical heating in air to the temperature of 1700°C. The ZrB2- SiC ceramic composites doped with a sintering additive in the Y2O3-Al2O3 system exhibited a high ablation resistance at the temperatures to 2800°C [39, 41].  It is highly important that the Y3Al5O12 distribution in the bulk should be maximally uniform, and its content should be minimized. This can be ensured by applying it as a coating on the surface of particles of the HfB2-30 vol % SiC composite powder by the sol-gel method.  The purpose of this work was to produce the HfB2- 30 vol % SiC composite powders modif ied with a small (≤ 5 vol %) amount of Y3Al5O12 and investigate their oxidation resistance.  EXPERIMENTAL  Reagents. Tetraethoxysilane (TEOS) Si(OC2H5)4 (>99.99%, AO EKOS-1), LBS-1 Bakelite varnish (OAO Karbolit), formic acid CH2O2 (>99%, OOO Spektr-Khim), and hafnium diboride (>98%, particle size 2-3 μm, aggregate size ~20-60 μm, OOO Tugoplavkie materialy). Metal acetylacetonates were  synthesized using Y(NO3)3 ⋅ 6H2O (99%, Khimmed), Al(NO3)3 ⋅ 9H2O (99%, Khimmed), C5H8O2 (>99.99%, AO EKOS-1), and 5% aqueous solution of NH3 · H2O (special-purity grade). The solvent of the obtained chelate coordination compounds and of the source of alkoxyl groups in the synthesis of heteroligand complexes was n-butanol C4H9OH (analytically pure). The IR transmission spectra of the solutions of the precursors after  the  replacement of  the C5H7O2 ligands were recorded with an InfraLUM FT-08 FTIR spectrometer (KBr glasses, 350-4000 cm-1). The X-ray powder diffraction patterns of the synthesized composite powders were recorded with a Bruker D8 ADVANCE X-ray powder diffractometer  radiation, resolution 0.02°, signal integration (CuKα time at point 0.3 s). The X-ray powder diffraction analysis was performed using “Match!”—Phase Identif ication from Powder Diffraction, Version 3.8.0.137 (Crystal Impact, Germany), into which the Crystallography Open Database (COD) is embedded. Scanning electron microscopy (SEM) studies were made with a Carl Zeiss NVision 40 focused ion beam scanning electron microscope. The elemental compositions of microregions were determined with an Oxford Instruments energy-dispersive X-ray analyzer. The thermal behavior of the synthesized products in an air f low (100 mL/min) in the temperature range 20-1200°C (heating rate 20 deg/min) was investigated with an SDT Q-600 simultaneous TGA/DSC/DTA analyzer.  RESULTS AND DISCUSSION  HfB2-SiC-Y3Al5O12 composite powders were produced  in two steps. At the f irst step, the sol-gel method [14, 62-64] was used to synthesize a HfB2- SiC powder in which the components were maximally uniformly distributed in each other. For this purpose, tetraethoxysilane (a silicon-containing precursor) and formic acid were introduced to a solution of phenol formaldehyde  resin, which was  further a carbon source. The molar ratios were n(Si) : n(CH2O2) = 1 : 4 and n(Si) : n(C) = 1 : 3.05. To the solution heated to ~50°C, water was added to a ratio of n(Si) : n(H2O) = 1 : 5. Immediately after the addition of water, a hafnium diboride powder was added (the calculated silicon carbide content of the HfB2-SiC system was 30 vol %), which was dispersed by simultaneous ultrasonication and mechanical stirring until gelation in the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', '1418  (a)  (c)  SIMONENKO et al.  (b)  2 µm  (d)  200 nm  1 µm  1 µm  Fig. 1. (a-c) SE (secondary electron) and (d) BSE (backscattered electron) SEM images of the microstructure of the HfB2- 30 vol % SiC composite powder.   system by the hydrolysis of tetraethoxysilane. The obtained gel was dried stepwise at temperatures of 50- 100°C, after which the xerogel was carbonized under dynamic vacuum conditions at 400°C for 2 h to form a reactive composition HfB2-(SiO2-C). The increased chemical activity of the nanostructured SiO2-C mixture enabled one to synthesize f inely divided silicon carbide at reduced pressure at 1550°C (5 h).  Modif ication  of  the HfB2-SiC  powder with Y3Al5O12 was performed by the sol-gel method with metal alkoxoacetylacetonates as precursors. The precursor [Al(C5H7O2)3 - x(C4H9On)x] was synthesized by heat treatment of a solution of aluminum acetylacetonate in an excess of 1-butanol in a round-bottom f lask with a ref lux condenser at an oil bath temperature of 160-165°C for 14 h. In this process, the chelate ligand was thermally decomposed to form acetone and butyl acetate (as was observed in the IR spectra of the solution), and the two substances were replaced by butoxo fragments. Further, yttrium acetylacetonate was added to the [Al(C5H7O2)3 - x(C4H9On)x] solution to a molar ratio of n(Al) : n(Y) = 5 : 3; the total content of metal cations was 0.2 mol/L. To complete the interligand exchange, the solution was stirred at ~50°C for 1 h.  To synthesize (HfB2-30 vol % SiC)-xY3Al5O12 (x = 2 and 5 vol %) composite powders, the HfB2-SiC powder was introduced to a required ratio to the solution of aluminum and yttrium alkoxoacetylacetonates and dispersed by ultrasonication and mechanical stirring. Then, the precursors were hydrolyzed by an ethanol ammonia solution as the system was heated to 60-70°C for 1 h, after which the samples were evapo rated until the weight ceased to change. The organic fragments of the Yand Al-containing xerogel were removed by heat treatment in air at 500°C for 2 h. The X-ray powder diffraction analysis of  the obtained powders showed that the synthesis of the f inely divided silicon carbide  is not accompanied either by oxidation of HfB2 particles (there are no ref lections of HfO2 phases), or by formation of hafnium carbide. The synthesized SiC is highly amorphized: its ref lections are not only weak (especially in comparison with the ref lections of the HfB2 phase), but also strongly broadened, resembling more the diffuse halos of partially ordered amorphous substances. The sintering additives Y3Al5O12 are X-ray amorphous, which was expected, taking into account the data [54, 65-67] that the crystallization temperature of the yttrium aluminum garnet phase in the sol-gel synthesis is >750-950°C (depending on the holding time). The scanning electron microscopy of the HfB2- 30 vol % SiC composite powder (Fig. 1) demonstrated that the produced silicon carbide is nanostructured. In addition, it was shown that the components are distributed between each other suff iciently uniformly: SiC agglomerates cover HfB2 particles ~1-2 μm in size or are located between them. The introduction of 2 vol % Y3Al5O12 leads to a certain change in the microstructure of the powder (Fig. 2); in particular, SiC particles coarsen, probably, by combining through an interlayer in the system Al2O3- Y2O3-SiO2. The surface of silicon carbide particles is known to almost always contain silicon dioxide, which  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1419  (а)  (c)  (b)  1 µm  (d)  1 µm  1 µm  200 nm  Fig. 2. (a, c, d) SE and (b) BSE SEM images of the microstructure of the HfB2-30 vol % SiC composite powder modif ied with 2 vol % Y3Al5O12.  is particularly signif icant for nanosized SiC particles. The fact that a thin layer of dielectric Y3Al5O12 coats not only silicon carbide agglomerates, but also the surface of hafnium diboride particles, was conf irmed by the phase-contrast images: the decrease in the quality of the microphotograph suggests an increase in the charge of the surface of the composite powders during SEM (Fig. 2b). Individual particles that could be assigned to Y3Al5O12 were not found. Still more considerable agglomeration of  f inely divided SiC particles was observed after addition of Y3Al5O12 to 5 vol % (Fig. 3). In the phase-contrast images (Fig. 3d), individual HfB2 or SiC particles are diff icult to discern, probably, because they are fully coated with a thin f ilm of yttrium aluminum oxide. Individual Y3Al5O12 particles were also not observed in the oxide-rich (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder.  The  thermal behavior of  the obtained (HfB2- 30 vol % SiC)-xY3Al5O12 (x = 0, 2, and 5 vol %) composite  powders  was  studied  by  simultaneous DSC/DTA/TGA in an air f low in the temperature range 20-1200°C (Fig. 4). As is seen, the oxidation processes at all the compositions occur at temperatures no lower than 600-650°C. The TGA curve of the HfB2-30 vol % SiC powder unmodif ied with Y3Al5O12 exhibits at least two steps of oxidation. The f irst step occurs at 710-780°C and is characterized by a sharp exothermic peak with a maximum at 756°C; at this step, there is obviously surface oxidation of HfB2. It is accompanied by an abrupt increase in the weight of the sample because of, f irst of all, the conversion of  HfB2 to HfO2 + B2O3. The higher-temperature exothermic events with maxima at 829 and 1035°C can be assigned to deeper oxidation of hafnium diboride that is complicated by the existence of a borosilicate glass layer  on  the  surface.  At  higher  temperatures (>1000°C), the additional decrease in the weight gain rate can be caused both by intensif ication of the oxidation of silicon carbide, which increases the SiO2 content of the molten glass on the surface and, correspondingly, its viscosity, and by more active evaporation from the surface of volatile boron oxide.  The  introduction  of  the  oxide  component (Y3Al5O12) to the composite powders gives rise to an additional exothermic event, a  shoulder at 695- 698°C, accompanied by an insignif icant (0.5-0.8%) weight loss and caused by the burnout of the residual carbon from the Aland Y-containing xerogel. The presence of such an amount of carbon has no signif icant effect on the oxidation resistance of the corresponding ceramic material obtained by hot pressing of the (HfB2-30 vol % SiC)-xY3Al5O12 powders, but guarantees the absence of HfO2 impurities on the surface of HfB2 grains. If the HfB2-30 vol % SiC powder is modif ied to only 2 vol %, the maximum of the exothermic event related to the primary oxidation of HfB2 is broadened and shifted toward higher temperatures by 23°C (from 756 to 779°C). Moreover, the shape of the weight gain curve also quite strongly changes (Fig. 4b): the weight 700-800°C gain  rate  in  the  temperature  range  decreases, which suggests the presence of a thin surface layer of amorphous Y3Al5O12, protecting from the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', '1420  (а)  (c)  SIMONENKO et al.  (b)  2 µm  (d)  Y Hf Si  Al  C  OHf  0 1  keV 2 3  4 µm  1 µm  1 µm  Fig. 3. (a-c) SE and (d) BSE SEM images of the microstructure of the HfB2-30 vol % SiC composite powder modif ied with 5 vol % Y3Al5O12.   interaction with oxygen. In the temperature range ~850-1050°C, the rates of the weight gain due to the oxidation for the unmodif ied powder and the powder containing 2 vol % Y3Al5O12 virtually coincide. At temperatures above 1050-1100°C, the oxidation rate insignif icantly decreases, probably, because of more active oxidation of SiC and the formation of the protective Al2O3-Y2O3-SiO2 layer rich in silicon dioxide. The total weight gain for this sample is 25.1% (from the minimum of the weight at 730°C), which almost coincides with that for the unmodif ied composite powder.  The increase in the Y3Al5O12 content to 5 vol % leads to a still larger shift of the temperature of the beginning of the oxidation of HfB2: the maximum of the most intense exothermic event is shifted by 52°C (from 756 to 808°C). The intensity of this exothermic event decreases  still  further, besides,  it broadens. A specif ic feature of the weight change curve in this case  is  that  the oxidation  rate  remains virtually unchanged throughout the temperature range 745- 1200°C. This can be explained by the fact that all the surface of HfB2 particles is coated with a continuous layer of amorphous oxide Y3Al5O12, and their oxidation rate is determined by the diffusion through the surface oxide layer. For the (HfB2-30 vol % SiC)- 5 vol % Y3Al5O12 composite, there are weak exothermic events at much higher temperatures (1056 and 1108°C) in comparison with the HfB2-30 vol % SiC  powder. Note that, for the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite, there is no signif icant change in the oxidation rate at temperatures >1080°C, which could be caused by the evaporation of B2O3. The total weight gain due to the oxidation of HfB2 and SiC is 24.8%.  CONCLUSIONS  ( ≤  The sol-gel method was used to synthesize the HfB2-30 vol % SiC composite powder in which f inely divided silicon carbide is maximally uniformly distributed between large (1-2-μm) hafnium diboride particles. The obtained powder was modif ied with a small 5 vol %) amount of Y3Al5O12 to increase the oxidation resistance and optimize the subsequent consolidation of ultra-high-temperature ceramic materials. As follows from the published data [41, 57] and the results of studying the evaporation of yttrium aluminum garnet at temperatures of 2200-2500°C [60], this oxide additive can also favor the stabilization of the hafnium dioxide, which forms by the oxidation of HfB2, in the cubic or tetragonal form; this should improve the adhesion of the oxidized layer during cooling or reheating of the ceramics. The controlled hydrolysis of the precursors (aluminum and yttrium alkoxoacetylacetonates) with subsequent gelation and annealing in air at 400°C produced the (HfB2-30 vol % SiC)-xY3Al5O12 (x = 2 and 5 vol %)  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1421  0%  2%  5%  carbon burnout, ~695°С  200  150  100  50  0 400  600  756°С  (а)  779°С  808°С  920°С  1056°С  1108°С  1035°С  1000  1200  829°С  800 T, °С  (b)  125  120  115  110  105  100  %  , t  h  g  i  e  W  108  106  10 4  102  100  700  800 T, °С  900  0%  2%  5%  W  m  ,  w  o  fl  t  a e  H  %  , t  h  g  i  e  W  400  600  800 T, °С  1000  1200  Fig. 4. (a) DSC and (b) TGA curves of the synthesized HfB2-30 vol % SiC composite powders: (black) unmodif ied and modif ied with (red) 2 and (blue) 5 vol % Y3Al5O12.  composite powders. The introduced X-ray amorphous phase Y3Al5O12 was distributed as a thin f ilm on particles of hafnium diboride and silicon carbide, leading to additional agglomeration of the latter. No formation of individual Y3Al5O12 particles was observed.  The investigation of the thermal behavior of the obtained composite powders in an air f low in the temperature range 20-1200°C showed that the modif ication to even 2 vol % Y3Al5O12 led to a signif icant decrease in the intensity of the low-temperature sur RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020          \\x0c', '1422  SIMONENKO et al.  face oxidation of HfB2 at 700-800°C: the weight gain rate decreased, and the corresponding exothermic peak decreased and broadened. In addition, the maximum of this exothermic peak was shifted toward higher temperatures. The total weight gain due to the oxidation was insignif icantly smaller than that for the unmodif ied HfB2-30 vol % SiC powder. The  introduction of 5 vol % Y3Al5O12 virtually completely neutralized the effect of the initial surface oxidation of HfB2: throughout the temperature range, the weight gain rate remained unchanged. Moreover, it shifted the temperature of the beginning of oxidation and the maximum of the corresponding thermal event from 756 to 808°C, and decreased the intensity of this thermal event. By and large, the introduction of 5 vol % Y3Al5O12 more signif icantly increased the oxidation resistance of the HfB2-30 vol % SiC composite, and the modifying oxide component was distributed maximally uniformly. The developed procedure of the synthesis of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder can be recommended for producing an ultrahigh-temperature ceramic of the corresponding composition.  FUNDING  This work was supported by the Russian Science Foundation (project no. 17-73-20181).  CONFLICT OF INTEREST  The authors declare that they have no conf licts of interest.  REFERENCES  1. 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Simonenko, A. N. Gordeev, et al., J. Eur. Ceram. Soc. 40, 1093 (2020).  https://doi.org/10.1016/j.jeurceramsoc.2019.11.023 65. N. P. Simonenko, E. P. Simonenko, V. G. Sevastyanov, and N. T. Kuznetsov, Russ. J. Inorg. Chem. 61, 667 (2016).  https://doi.org/10.1134/S003602361606019X 66. E. P. Simonenko, N. P. Simonenko, V. G. Sevastyanov, and N. T. Kuznetsov, Russ. J. Inorg. Chem. 57, 1521 (2012).  https://doi.org/10.1134/S0036023612120194 67. E. P. Simonenko, N. P. Simonenko, G. P. Kopitsa, et al., Russ. J. Inorg. Chem. 63, 691 (2018).  https://doi.org/10.1134/S0036023618060232  Translated by V. Glyanchenko  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c']"
},{
  "_id": 224,
  "PDF": "Production-and-Oxidation-Resistance-of-HfB30-vol--SiC-Composite-Powders-Modified-with-YAlO2020Russian-Journal-of-Inorganic-Chemistry.pdf",
  "Text": "['ISSN 0036-0236, Russian Journal of Inorganic Chemistry, 2020, Vol. 65, No. 9, pp. 1416-1423. © Pleiades Publishing, Ltd., 2020. Russian Text © The Author(s), 2020, published in Zhurnal Neorganicheskoi Khimii, 2020, Vol. 65, No. 9, pp. 1274-1282.  INORGANIC MATERIALS  AND NANOMATERIALS  Production and Oxidation Resistance of HfB2-30 vol % SiC  Composite Powders Modif ied with Y3Al5O12  E. P. Simonenkoa, *, N. P. Simonenkoa, I. A . Nagornova, b, V. G. Sevastyanova, and N. T. Kuznetsova  aKurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences, Moscow, 119991 Russia bMendeleev University of Chemical Technology of Russia, Moscow, 125047 Russia *e-mail: ep_simonenko@mail.ru  Received March 16, 2020; revised April 13, 2020; accepted April 27, 2020  Abstract—The HfB2-30 vol % SiC composite powders modif ied with a small (2 and 5 vol %) amount of Y3Al5O12 were synthesized by the sol-gel method. Such a modif ication in producing an ultra-high-temperature ceramic of the corresponding composition optimized the process temperature and increased the oxidation resistance of the ceramic by stabilizing the formed hafnium dioxide in the cubic or tetragonal form. It was determined that the introduced X-ray amorphous phase Y3Al5O12 was distributed as a thin f ilm on particles of hafnium diboride and silicon carbide and led to agglomeration of the latter; no formation of individual Y3Al5O12 particles was observed. Investigation of the thermal behavior of the obtained composite powders in an air f low in the temperature range 20-1200°C showed that the introduction of 5 vol % Y3Al5O12 virtually completely neutralized the effect of the initial surface oxidation of HfB2 in the temperature range 700-800°C, shifted the temperature of the beginning of oxidation and the maximum of the corresponding thermal event from 756 to 808°C, and decreased the intensity of this thermal event.  Keywords: UHTC, composite powder, HfB2, SiC, YAG, sol-gel method, oxidation resistance DOI: 10.1134/S003602362009020X  INTRODUCTION  Numerous  research  groups develop MB2-SiC (M = Zr, Hf) materials proposed to be used during aerodynamic heating to temperatures about ~2500°C [1-15]. Along with a suff icient resistance to oxidation, including oxidation by atomic oxygen [16-20], such ceramic composites have reasonably good mechanical performance, high thermal conductivity (also at elevated temperatures), optical properties suitable for potential applications, and low catalytic activity in exothermic reactions of surface recombination of oxygen and nitrogen atoms, etc. One of  the key properties of ZrB2(HfB2)-SiC ultra-high-temperature ceramics (UHTC) is the ability to endure quite a long exposure to an oxygen-containing atmosphere at elevated temperatures (>1200- 2500°C) owing to the presence of a protective layer of borosilicate glass  formed by  the oxidation of  the ceramic. Depending on  the characteristics of  the material (composition, porosity) and the exposure conditions, this airproof glassy layer can occur both on the surface of the ceramic (at not too high temperatures of the latter: <1700-1850°C), and in the bulk of the porous layer based on the hafnium oxide, a product of the oxidation of HfB2 (at high surface temperature >1900-2000°C). Therefore, investigation of the oxidation resistance of UHTC of a specif ic composi tion (with taking into account various strengthening additives) and the microstructure under certain exposure conditions are extremely important [21-38]. One of the obvious methods to increase the oxidation resistance is to develop methods to obtain maximally dense materials. The structures of zirconium and hafnium diborides and silicon carbide are such that  the production of ceramics based on  them requires quite high temperatures, which leads not only to high energy consumption, but also to impairment of the mechanical performance because of coarsening of grains (f irst of all, of metal diboride) by high-temperature consolidation. Introduction of sintering additives signif icantly reduces the required process temperature. Liquid-phase sintering of oxygen-free ceramics using additives in the system Al2O3-Y2O3, including an additive corresponding to yttrium aluminum garnet Y3Al5O12 (YAG), has long been used to ensure excellent densif ication at not too high temperatures both in pressureless sintering [39-44], and in high-temperature pressing: hot pressing [45-49] or spark plasma sintering [50-54]. The amount, particle size, and the uniformity of the distribution of a sintering additive affect signif icantly not only the densif ication process, but also  the properties of  the obtained material, including the oxidation resistance. However, it was  1416  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1417  noted that a high YAG content in long-term treatment at elevated temperatures and low pressure can lead to degradation of the material [48-56] because of the interaction of the components of the ceramic. Therefore, a most important problem is to minimize the necessary amount of an introduced sintering additive.  One  more  problem  of  the  oxidation  of ZrB2(HfB2)-SiC UHTC by a high-enthalpy air f low is the separation of the oxidized layer because of phase transformations during cooling or reheating of such oxidation products as high-melting ZrO2/HfO2, the thermodynamically stable form of which is the monoclinic crystal lattice. A solution can be modif ication of the ceramic with agents capable of stabilizing hafnium oxide in the cubic or tetragonal system, e.g., with high-melting and low-volatile yttrium oxide [57-59]. For example, it was shown [58] that, in the HfB2- 20 vol % SiC-3 vol % Y2O3 composite (in comparison with the unmodif ied material), a protective surface layer of borosilicate glass is stabilized.  However, a technologically better stabilizing additive could be less high-melting yttrium aluminum garnet. It was demonstrated [60] that, at a temperature >2200°C, from Y3Al5O12, more volatile Al2O3 can be distilled off, and the remaining Y2O3 can be consumed for the stabilization of the product of the oxidation of HfB2. At a lower temperature, the unevaporated alumina  can modify  the  silicate  glass,  specif ically, increase its softening temperatures and the viscosity in the molten state, thus decreasing the oxygen diffusion into the bulk of the material ZrB2(HfB2)-SiC. The eff iciency of the introduction of Y3Al5O12 for improving the oxygen resistance of UHTC has recently been conf irmed [46]: it was concluded that the doping of the ZrB2-SiC ceramics with Al/Y oxides considerably improved their oxidation resistance at the temperature of 1700°C because of a decrease in the activity of silicon dioxide. The ZrB2-SiC-B4C-Y3Al5O12 coatings [61] had a good thermal-shock resistance in cyclical heating in air to the temperature of 1700°C. The ZrB2- SiC ceramic composites doped with a sintering additive in the Y2O3-Al2O3 system exhibited a high ablation resistance at the temperatures to 2800°C [39, 41].  It is highly important that the Y3Al5O12 distribution in the bulk should be maximally uniform, and its content should be minimized. This can be ensured by applying it as a coating on the surface of particles of the HfB2-30 vol % SiC composite powder by the sol-gel method.  The purpose of this work was to produce the HfB2- 30 vol % SiC composite powders modif ied with a small (≤ 5 vol %) amount of Y3Al5O12 and investigate their oxidation resistance.  EXPERIMENTAL  Reagents. Tetraethoxysilane (TEOS) Si(OC2H5)4 (>99.99%, AO EKOS-1), LBS-1 Bakelite varnish (OAO Karbolit), formic acid CH2O2 (>99%, OOO Spektr-Khim), and hafnium diboride (>98%, particle size 2-3 μm, aggregate size ~20-60 μm, OOO Tugoplavkie materialy). Metal acetylacetonates were  synthesized using Y(NO3)3 ⋅ 6H2O (99%, Khimmed), Al(NO3)3 ⋅ 9H2O (99%, Khimmed), C5H8O2 (>99.99%, AO EKOS-1), and 5% aqueous solution of NH3 · H2O (special-purity grade). The solvent of the obtained chelate coordination compounds and of the source of alkoxyl groups in the synthesis of heteroligand complexes was n-butanol C4H9OH (analytically pure). The IR transmission spectra of the solutions of the precursors after  the  replacement of  the C5H7O2 ligands were recorded with an InfraLUM FT-08 FTIR spectrometer (KBr glasses, 350-4000 cm-1). The X-ray powder diffraction patterns of the synthesized composite powders were recorded with a Bruker D8 ADVANCE X-ray powder diffractometer  radiation, resolution 0.02°, signal integration (CuKα time at point 0.3 s). The X-ray powder diffraction analysis was performed using “Match!”—Phase Identif ication from Powder Diffraction, Version 3.8.0.137 (Crystal Impact, Germany), into which the Crystallography Open Database (COD) is embedded. Scanning electron microscopy (SEM) studies were made with a Carl Zeiss NVision 40 focused ion beam scanning electron microscope. The elemental compositions of microregions were determined with an Oxford Instruments energy-dispersive X-ray analyzer. The thermal behavior of the synthesized products in an air f low (100 mL/min) in the temperature range 20-1200°C (heating rate 20 deg/min) was investigated with an SDT Q-600 simultaneous TGA/DSC/DTA analyzer.  RESULTS AND DISCUSSION  HfB2-SiC-Y3Al5O12 composite powders were produced  in two steps. At the f irst step, the sol-gel method [14, 62-64] was used to synthesize a HfB2- SiC powder in which the components were maximally uniformly distributed in each other. For this purpose, tetraethoxysilane (a silicon-containing precursor) and formic acid were introduced to a solution of phenol formaldehyde  resin, which was  further a carbon source. The molar ratios were n(Si) : n(CH2O2) = 1 : 4 and n(Si) : n(C) = 1 : 3.05. To the solution heated to ~50°C, water was added to a ratio of n(Si) : n(H2O) = 1 : 5. Immediately after the addition of water, a hafnium diboride powder was added (the calculated silicon carbide content of the HfB2-SiC system was 30 vol %), which was dispersed by simultaneous ultrasonication and mechanical stirring until gelation in the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', '1418  (a)  (c)  SIMONENKO et al.  (b)  2 µm  (d)  200 nm  1 µm  1 µm  Fig. 1. (a-c) SE (secondary electron) and (d) BSE (backscattered electron) SEM images of the microstructure of the HfB2- 30 vol % SiC composite powder.   system by the hydrolysis of tetraethoxysilane. The obtained gel was dried stepwise at temperatures of 50- 100°C, after which the xerogel was carbonized under dynamic vacuum conditions at 400°C for 2 h to form a reactive composition HfB2-(SiO2-C). The increased chemical activity of the nanostructured SiO2-C mixture enabled one to synthesize f inely divided silicon carbide at reduced pressure at 1550°C (5 h).  Modif ication  of  the HfB2-SiC  powder with Y3Al5O12 was performed by the sol-gel method with metal alkoxoacetylacetonates as precursors. The precursor [Al(C5H7O2)3 - x(C4H9On)x] was synthesized by heat treatment of a solution of aluminum acetylacetonate in an excess of 1-butanol in a round-bottom f lask with a ref lux condenser at an oil bath temperature of 160-165°C for 14 h. In this process, the chelate ligand was thermally decomposed to form acetone and butyl acetate (as was observed in the IR spectra of the solution), and the two substances were replaced by butoxo fragments. Further, yttrium acetylacetonate was added to the [Al(C5H7O2)3 - x(C4H9On)x] solution to a molar ratio of n(Al) : n(Y) = 5 : 3; the total content of metal cations was 0.2 mol/L. To complete the interligand exchange, the solution was stirred at ~50°C for 1 h.  To synthesize (HfB2-30 vol % SiC)-xY3Al5O12 (x = 2 and 5 vol %) composite powders, the HfB2-SiC powder was introduced to a required ratio to the solution of aluminum and yttrium alkoxoacetylacetonates and dispersed by ultrasonication and mechanical stirring. Then, the precursors were hydrolyzed by an ethanol ammonia solution as the system was heated to 60-70°C for 1 h, after which the samples were evapo rated until the weight ceased to change. The organic fragments of the Yand Al-containing xerogel were removed by heat treatment in air at 500°C for 2 h. The X-ray powder diffraction analysis of  the obtained powders showed that the synthesis of the f inely divided silicon carbide  is not accompanied either by oxidation of HfB2 particles (there are no ref lections of HfO2 phases), or by formation of hafnium carbide. The synthesized SiC is highly amorphized: its ref lections are not only weak (especially in comparison with the ref lections of the HfB2 phase), but also strongly broadened, resembling more the diffuse halos of partially ordered amorphous substances. The sintering additives Y3Al5O12 are X-ray amorphous, which was expected, taking into account the data [54, 65-67] that the crystallization temperature of the yttrium aluminum garnet phase in the sol-gel synthesis is >750-950°C (depending on the holding time). The scanning electron microscopy of the HfB2- 30 vol % SiC composite powder (Fig. 1) demonstrated that the produced silicon carbide is nanostructured. In addition, it was shown that the components are distributed between each other suff iciently uniformly: SiC agglomerates cover HfB2 particles ~1-2 μm in size or are located between them. The introduction of 2 vol % Y3Al5O12 leads to a certain change in the microstructure of the powder (Fig. 2); in particular, SiC particles coarsen, probably, by combining through an interlayer in the system Al2O3- Y2O3-SiO2. The surface of silicon carbide particles is known to almost always contain silicon dioxide, which  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1419  (а)  (c)  (b)  1 µm  (d)  1 µm  1 µm  200 nm  Fig. 2. (a, c, d) SE and (b) BSE SEM images of the microstructure of the HfB2-30 vol % SiC composite powder modif ied with 2 vol % Y3Al5O12.  is particularly signif icant for nanosized SiC particles. The fact that a thin layer of dielectric Y3Al5O12 coats not only silicon carbide agglomerates, but also the surface of hafnium diboride particles, was conf irmed by the phase-contrast images: the decrease in the quality of the microphotograph suggests an increase in the charge of the surface of the composite powders during SEM (Fig. 2b). Individual particles that could be assigned to Y3Al5O12 were not found. Still more considerable agglomeration of  f inely divided SiC particles was observed after addition of Y3Al5O12 to 5 vol % (Fig. 3). In the phase-contrast images (Fig. 3d), individual HfB2 or SiC particles are diff icult to discern, probably, because they are fully coated with a thin f ilm of yttrium aluminum oxide. Individual Y3Al5O12 particles were also not observed in the oxide-rich (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder.  The  thermal behavior of  the obtained (HfB2- 30 vol % SiC)-xY3Al5O12 (x = 0, 2, and 5 vol %) composite  powders  was  studied  by  simultaneous DSC/DTA/TGA in an air f low in the temperature range 20-1200°C (Fig. 4). As is seen, the oxidation processes at all the compositions occur at temperatures no lower than 600-650°C. The TGA curve of the HfB2-30 vol % SiC powder unmodif ied with Y3Al5O12 exhibits at least two steps of oxidation. The f irst step occurs at 710-780°C and is characterized by a sharp exothermic peak with a maximum at 756°C; at this step, there is obviously surface oxidation of HfB2. It is accompanied by an abrupt increase in the weight of the sample because of, f irst of all, the conversion of  HfB2 to HfO2 + B2O3. The higher-temperature exothermic events with maxima at 829 and 1035°C can be assigned to deeper oxidation of hafnium diboride that is complicated by the existence of a borosilicate glass layer  on  the  surface.  At  higher  temperatures (>1000°C), the additional decrease in the weight gain rate can be caused both by intensif ication of the oxidation of silicon carbide, which increases the SiO2 content of the molten glass on the surface and, correspondingly, its viscosity, and by more active evaporation from the surface of volatile boron oxide.  The  introduction  of  the  oxide  component (Y3Al5O12) to the composite powders gives rise to an additional exothermic event, a  shoulder at 695- 698°C, accompanied by an insignif icant (0.5-0.8%) weight loss and caused by the burnout of the residual carbon from the Aland Y-containing xerogel. The presence of such an amount of carbon has no signif icant effect on the oxidation resistance of the corresponding ceramic material obtained by hot pressing of the (HfB2-30 vol % SiC)-xY3Al5O12 powders, but guarantees the absence of HfO2 impurities on the surface of HfB2 grains. If the HfB2-30 vol % SiC powder is modif ied to only 2 vol %, the maximum of the exothermic event related to the primary oxidation of HfB2 is broadened and shifted toward higher temperatures by 23°C (from 756 to 779°C). Moreover, the shape of the weight gain curve also quite strongly changes (Fig. 4b): the weight 700-800°C gain  rate  in  the  temperature  range  decreases, which suggests the presence of a thin surface layer of amorphous Y3Al5O12, protecting from the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', '1420  (а)  (c)  SIMONENKO et al.  (b)  2 µm  (d)  Y Hf Si  Al  C  OHf  0 1  keV 2 3  4 µm  1 µm  1 µm  Fig. 3. (a-c) SE and (d) BSE SEM images of the microstructure of the HfB2-30 vol % SiC composite powder modif ied with 5 vol % Y3Al5O12.   interaction with oxygen. In the temperature range ~850-1050°C, the rates of the weight gain due to the oxidation for the unmodif ied powder and the powder containing 2 vol % Y3Al5O12 virtually coincide. At temperatures above 1050-1100°C, the oxidation rate insignif icantly decreases, probably, because of more active oxidation of SiC and the formation of the protective Al2O3-Y2O3-SiO2 layer rich in silicon dioxide. The total weight gain for this sample is 25.1% (from the minimum of the weight at 730°C), which almost coincides with that for the unmodif ied composite powder.  The increase in the Y3Al5O12 content to 5 vol % leads to a still larger shift of the temperature of the beginning of the oxidation of HfB2: the maximum of the most intense exothermic event is shifted by 52°C (from 756 to 808°C). The intensity of this exothermic event decreases  still  further, besides,  it broadens. A specif ic feature of the weight change curve in this case  is  that  the oxidation  rate  remains virtually unchanged throughout the temperature range 745- 1200°C. This can be explained by the fact that all the surface of HfB2 particles is coated with a continuous layer of amorphous oxide Y3Al5O12, and their oxidation rate is determined by the diffusion through the surface oxide layer. For the (HfB2-30 vol % SiC)- 5 vol % Y3Al5O12 composite, there are weak exothermic events at much higher temperatures (1056 and 1108°C) in comparison with the HfB2-30 vol % SiC  powder. Note that, for the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite, there is no signif icant change in the oxidation rate at temperatures >1080°C, which could be caused by the evaporation of B2O3. The total weight gain due to the oxidation of HfB2 and SiC is 24.8%.  CONCLUSIONS  ( ≤  The sol-gel method was used to synthesize the HfB2-30 vol % SiC composite powder in which f inely divided silicon carbide is maximally uniformly distributed between large (1-2-μm) hafnium diboride particles. The obtained powder was modif ied with a small 5 vol %) amount of Y3Al5O12 to increase the oxidation resistance and optimize the subsequent consolidation of ultra-high-temperature ceramic materials. As follows from the published data [41, 57] and the results of studying the evaporation of yttrium aluminum garnet at temperatures of 2200-2500°C [60], this oxide additive can also favor the stabilization of the hafnium dioxide, which forms by the oxidation of HfB2, in the cubic or tetragonal form; this should improve the adhesion of the oxidized layer during cooling or reheating of the ceramics. The controlled hydrolysis of the precursors (aluminum and yttrium alkoxoacetylacetonates) with subsequent gelation and annealing in air at 400°C produced the (HfB2-30 vol % SiC)-xY3Al5O12 (x = 2 and 5 vol %)  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c', 'PRODUCTION AND OXIDATION RESISTANCE OF HfB2-30 VOL % SiC  1421  0%  2%  5%  carbon burnout, ~695°С  200  150  100  50  0 400  600  756°С  (а)  779°С  808°С  920°С  1056°С  1108°С  1035°С  1000  1200  829°С  800 T, °С  (b)  125  120  115  110  105  100  %  , t  h  g  i  e  W  108  106  10 4  102  100  700  800 T, °С  900  0%  2%  5%  W  m  ,  w  o  fl  t  a e  H  %  , t  h  g  i  e  W  400  600  800 T, °С  1000  1200  Fig. 4. (a) DSC and (b) TGA curves of the synthesized HfB2-30 vol % SiC composite powders: (black) unmodif ied and modif ied with (red) 2 and (blue) 5 vol % Y3Al5O12.  composite powders. The introduced X-ray amorphous phase Y3Al5O12 was distributed as a thin f ilm on particles of hafnium diboride and silicon carbide, leading to additional agglomeration of the latter. No formation of individual Y3Al5O12 particles was observed.  The investigation of the thermal behavior of the obtained composite powders in an air f low in the temperature range 20-1200°C showed that the modif ication to even 2 vol % Y3Al5O12 led to a signif icant decrease in the intensity of the low-temperature sur RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020          \\x0c', '1422  SIMONENKO et al.  face oxidation of HfB2 at 700-800°C: the weight gain rate decreased, and the corresponding exothermic peak decreased and broadened. In addition, the maximum of this exothermic peak was shifted toward higher temperatures. The total weight gain due to the oxidation was insignif icantly smaller than that for the unmodif ied HfB2-30 vol % SiC powder. The  introduction of 5 vol % Y3Al5O12 virtually completely neutralized the effect of the initial surface oxidation of HfB2: throughout the temperature range, the weight gain rate remained unchanged. Moreover, it shifted the temperature of the beginning of oxidation and the maximum of the corresponding thermal event from 756 to 808°C, and decreased the intensity of this thermal event. By and large, the introduction of 5 vol % Y3Al5O12 more signif icantly increased the oxidation resistance of the HfB2-30 vol % SiC composite, and the modifying oxide component was distributed maximally uniformly. The developed procedure of the synthesis of the (HfB2-30 vol % SiC)-5 vol % Y3Al5O12 composite powder can be recommended for producing an ultrahigh-temperature ceramic of the corresponding composition.  FUNDING  This work was supported by the Russian Science Foundation (project no. 17-73-20181).  CONFLICT OF INTEREST  The authors declare that they have no conf licts of interest.  REFERENCES  1. 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Simonenko, A. N. Gordeev, et al., J. Eur. Ceram. Soc. 40, 1093 (2020).  https://doi.org/10.1016/j.jeurceramsoc.2019.11.023 65. N. P. Simonenko, E. P. Simonenko, V. G. Sevastyanov, and N. T. Kuznetsov, Russ. J. Inorg. Chem. 61, 667 (2016).  https://doi.org/10.1134/S003602361606019X 66. E. P. Simonenko, N. P. Simonenko, V. G. Sevastyanov, and N. T. Kuznetsov, Russ. J. Inorg. Chem. 57, 1521 (2012).  https://doi.org/10.1134/S0036023612120194 67. E. P. Simonenko, N. P. Simonenko, G. P. Kopitsa, et al., Russ. J. Inorg. Chem. 63, 691 (2018).  https://doi.org/10.1134/S0036023618060232  Translated by V. Glyanchenko  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY    Vol. 65    No. 9    2020  \\x0c']"
},{
  "_id": 225,
  "PDF": "Progress in the fabrication of ultra-high-temperature ceramics in situ synthesis, microstructure and properties of a reactive hot-pressed HfB2–SiC composite.pdf",
  "Text": "['Composites Science and Technology 65 (2005) 1869-1879  COMPOSITES SCIENCE AND TECHNOLOGY  www.elsevier.com/locate/compscitech  Progress in the fabrication of ultra-high-temperature ceramics:  ‘‘in situ’’ synthesis, microstructure and properties of a reactive  hot-pressed HfB2-SiC composite  Fre´de´ ric Monteverde *  National Research Council, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy  Received 4 October 2004; received in revised form 23 March 2005; accepted 9 April 2005  Available online 12 May 2005  Abstract  An ultra-high-temperature HfB2-SiC composite was successfully fabricated by reactive hot-pressing. Solid reagents like Hf/Si/ B4C, mechanically mixed in molar ratio 2.2/0.8/1, were ‘‘in situ’’ converted into the basic ingredients (i.e., HfB2, SiC), and then directly hot-pressed until full density was achieved (1900 °C final temperature). The microstructure consisted of faceted diboride grains (mean size 3 lm), with HfC (6 vol%) and SiC (22 vol% and mean size 1 lm) evenly distributed intergranularly. The combi nation of some mechanical properties was of considerable significance: about 19 GPa of micro-hardness, 520 GPa of YoungÕs modulus, 770 ± 35 and 315 ± 10 MPa of flexural strength at 25 and 1500 °C, respectively. A relevant merit characterized the resistance to oxidation: repeated exposures at 1700 °C, or at 1450 °C for 20 h,  involved limited mass gains and small changes of the original  microstructure. The marked refractoriness of HfB2 and SiC, which constitute the framework of cially its thermostructural stability.  Ó 2005 Elsevier Ltd. All rights reserved.  the composite, dominates benefi Keywords: Ultra-high temperature ceramic (UHTCs); Microstructure; Thermomechanical properties; Oxidation; Reactive hot-pressing  1. Introduction  The design and processing of materials with enhanced  ultra-high-temperature ceramics (UHTCs) capable to in crease the toleration of heat on sharp leading profiles of space vehicles reentering the EarthÕs atmosphere [2-5].  high temperature capabilities represent one of  the most  During recent decades, a core of attention of the sci challenging  task  of modern  engineering. Within  the  entific  community has been focused on the design of  crowded families of advanced engineered ceramics, tran ceramic matrix composites (CMCs), owing to the shared  sition metal diborides and carbides are naturally selected  opinion  among  researchers  that  the  development  of  for ultra-high-temperature  structural  applications  (for  monolithic ceramics may secure only marginal perfor instance  furnace  elements,  crucibles,  arc-plasma  elec mance benefits. CMCs have been largely produced by  trodes, thermal protection) because of melting points exceeding 3000 °C, coupled to an overall  thermostruc densification of mechanically mixed powders. Since the  melting points of  the UHTCs are among the highest  tural stability at very high temperatures [1]. Great atten known, processing them into CMCs with final complex  tion is currently addressed towards  the engineering of  shapes and full density requires long exposure in atmo * Tel.: +390546699758; fax: +39054646381.  E-mail address: fmonte@istec.cnr.it.  0266-3538/$ see front matter Ó 2005 Elsevier Ltd. All rights reserved.  doi:10.1016/j.compscitech.2005.04.003  sphere-controlled  conventional  furnaces  at  extremely  high sintering temperatures and applied pressures.  As  an alternative  route, CMCs have been proved  being manufactured using  ‘‘in situ’’ high-temperature  \\x0c', 'solid-state  chemical displacement  reactions  [6-8]. The  advantage of this approach is that it has offered the pos sibility to obtain composites not only with controlled  microstructure, but also characterized by an high chem ical  compatibility of  the  ‘‘in situ’’  formed individual  phases  evenly  distributed. With  the  ability  to  adjust  the final microstructure being understood,  the proposal  of unexplored routes  is additionally motivated by the  demand for reducing manufacturing costs (for instance,  cheaper raw powders and processing steps). Unlike self propagating  high-temperature  synthesis  (SHS),  that  basically  exploits  the  exothermicity  of  uncontrolled  chemical  reactions,  a  displacement  reaction  synthesis  can be accomplished gradually via solid-state diffusion  at  temperatures below that  ignited once a (triggered)  SHS takes place [9]. It ensures better control over the fi nal microstructure and the isotropy of properties.  Superior UHTCs would be a great asset to improving  the  capabilities  of  spacecrafts. Current  research  into  thermal  insulating structures is aimed at designing inno vative UHTCs  for  sharp leading edges of hypersonic  space  vehicles. These  new materials  should  combine  the potentiality to withstand extremely high tempera tures to the capacity of dissipating heat efficiently [1,3].  The poor  sinterability and the resistance to oxidation/  ablation, which  represent  topical  deficiencies  in  the  UHTC class of materials, are addressed with an appro priate  focus on compositional designs  and composite  constructions [3,10,11].  Following what has been just  reported,  the present  paper  intended to produce by reactive hot-pressing an  HfB2-SiC composite which (to date) does not draw on already published data. ‘‘In situ’’ synthesis and sinter ing, microstructure and some thermostructural proper ties were studied and discussed.  2. Experimental  Three commercial solid precursors like Hf pure, FSSS 2.1 lm, Neomat Co. - Latvia,), Si (grade AX05 99.9% pure, FSSS 3.5 lm, H.C. Starck - Germany,) and B4C (grade HS, FSSS 0.8 lm, H.C. Starck - Germany) were selected. In accordance with the fol (99.8%  lowing reaction  ð2 þ xÞHf þ ð1 \\x00 xÞSi þ B4C  ) 2HfB2 þ ð1 \\x00 xÞSiC þ xHfC  ð1Þ  the  correspondent  stoichiometric  precursors  were  weighed for x = 0.2. Such a formulation, adjusted by  adding limited quantities of HfB2 (99.5% pure, 325 mesh, Cerac Inc. - USA) and a-SiC (grade UF25, FSSS 0.45 lm, H.C. Starck - Germany), was milled for 24 h in  a  polyethylene  jar  using  absolute  ethanol  and ZrO2 balls, dried with a rotary evaporator under a continuous  stream of nitrogen, and sieved. The expected final com position (vol%) reads HfB2 + 22.1 SiC + 5.9 HfC. In order to investigate the end-product formation  kinetics of  reaction (1), a rational  campaign of  (pres sureless) heat  treatments, hereafter  labelled PLSHT-n,  was conducted on the as-ground powder mix from 1000 °C up to 1650 °C in flowing Argon, with a heating rate of 10 °C/min and 1 h of dwell  time, using graphite  heating elements and crucibles (Astro Industries Inc. -  USA). Compacted pellets  (about 2 g in weight, 12 mm  in diameter) were isostatically 3500 kg cm\\x002. After the heat  cold-pressed  at  treatments,  the PLSHT-n  pellets were finely crushed in an agate mortar, and then analyzed via X-ray diffraction (XRD, Ni-filtered Cu Ka  radiation, mod. D500, Siemens - Germany). Thermody namic calculations were performed using the HSC soft ware package [12].  The reactive hot-pressing was performed at  low vac uum (0.5 mbar)  using  a BN-lined  induction-heated  graphite  die,  into which  the  as-ground  powder mix  was directly loaded and heat-treated. The  schedule of  the thermal treatment is depicted in Fig. 1. The final set point of the hot-press run was 1900 °C. The temper ature was measured with a pyrometer  focused on the  graphite die.  The bulk density was measured using the Archime dian method, while  the  relative density was  estimated  applying the  rule of mixture. The microstructure was  analyzed with a  scanning  electron microscope  (SEM,  Leica Cambridge, mod. S360, UK) equipped with an en ergy dispersive microanalyzer  (EDX, mod.  INCA En ergy 300, Oxford Instruments Analytical, UK), and an  X-ray diffractometer as well. Polished sections (finished 0.5 lm) of  the  reactive hot-pressed material were pre pared using diamond abrasives.  Micro-hardness  (Hv1.0) was measured by a Vickers  indenter with 9.81 N as applied load for 15 s on a polished section. Flexural strength (r) in a 4-pt. configuration was tested at two temperatures, 25 and 1500 °C, ambient air on 25.0 · 2.5 · 2.0 mm3 (length · width · thickness,  in  chamfered  bar  respectively), using 20 and  10 mm as outer and inner span, cross-head speed of 0.5 mm min\\x001.  respectively,  and  a  Likewise,  two  oxidation  treatments  (at  ambient  pressure)  (T-1)  isothermal run at 1450 °C for 20 h in flowing dry (50 cm3 min\\x001), 30 °C min\\x001 of heating rate  air  and free cooling; isothermal cycled run at 1700 °C, 10 min of expo (T-2)  sure each,  loading and removal of the coupon at  the fixed set-point  were carried out. 2.5 · 2.0 · 8.0 mm3  Coupons with dimensions of finish Ra \\x18 0.2 lm) were washed in an ultrasonic bath of acetone, and then dried  (surface  1870  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  \\x0c', 'at 80 °C overnight. Treatment T-1 was executed using a  thermogravimetric analyzer (mod. STA409, NETZSCH Gera¨ tebau GmbH - Germany), 10\\x003 mg of accuracy,  equipped with a vertically heated Al2O3 chamber. The test piece was placed upon zirconia supports, separating  them from the Al2O3 holder. Treatment T-2 was performed using a bottom-loading box furnace, heated with  MoSi2 elements. The test specimen was placed upon SiC supports. The mass of the sample was measured before  and after each cycle. Polished cross-sections of  the oxi dized samples were prepared and analyzed via SEM-  EDX. The  specimens  (polished)  surface,  imaged using  secondary electrons (SEs), were all observed free of con ductive coating in order to maintain the sensitivity to the  light elements of  the EDX equipment as high as possi ble.  In addition, using an optical microscope,  the final  thickness of the oxidized samples was measured as well.  3. Results  3.1. The ceramic synthesis via solid-state precursors  Processing parameters and some results from the ser ies of  the pressureless heat  treatments  (PLSHT-n) are  shown in Table 1. Stacked XRD patterns from the start ing as-ground mixture (PLSHT-0) and after treatments at 1000, 1100 and 1200 °C are presented (Fig. 2). In from a-Hf, Si and B4C, also a cubic modified  the mixture PLSHT-0, apart  X-ray  diffraction  identified  HfH1.5 (ICDD 05-639) and HfB2. The former was tially present in the as-received hafnium powder because  ini of the wet storage. Such a hafnium hydride proved to be stable up to 1100 °C.  Instead,  the  formation of HfB2 during ball-milling clearly highlights how reaction (1) succeeded (\\x00695 kJ/mol  is prone  to have  enthalpy of  1530 1500  1660  1820  1450 x 60 min  1900  50  0.5  10  19  34  39  1400  1500  1600  1700  1800  1900  2000  80  100  120  140  160  180  200  Time / min  T  /  ˚  C  0  10  20  30  40  50  P  /  M  P  a  Temperature  Pressure  Fig. 1. Schedule of the reactive hot-pressing experiment: temperature (T, temperature of 1450 °C was reached in about 90 min.  left Y-axis) and applied pressure (P, right Y-axis) vs. processing time. The  Table 1  Parameters and results of the PLSHT-n tests: temperature T, dwell  time t, relative weight loss WL, and crystalline end-products  Test  T (°C)  t (min)  WL (%)  Crystalline end-productsb  Main  Minor  PLSHT-1  1000  60  1.3  Hfa, Si, HfB2, HfC, HfH1.5, B4 C Hfa, (Hf,Si), HfB2, HfC  (Hf,Si)  PLSHT-2  1100  60  1.5  -  PLSHT-3  1200  60  2.3  HfB2, SiC, HfC  m-HfO2  PLSHT-4  1300  60  1.0  HfB2, SiC, HfC  m-HfO2  PLSHT-5  1450  60  0.5  HfB2, SiC, HfC  m-HfO2  PLSHT-6  1500  60  1.0  HfB2, SiC, HfC  m-HfO2  PLSHT-7  1500  120  1.1  HfB2, SiC, HfC  m-HfO2  PLSHT-8  1550  60  1.4  HfB2, SiC, HfC  m-HfO2  c  PLSHT-9  1600  60  1.2  HfB2, SiC, HfC  -  PLSHT-10  1650  60  1.5  HfB2, SiC, HfC  -  a Hexagonal type (ICDD 38-1478) but longer lattice parameters. b Pulverized sample. c Uncertain.  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  1871    \\x0c', 'formation at 25 °C),  and satisfies also the  thermody namic conditions for a self-sustaining reaction process (TADIAB ffi 2900 °C). In accordance with the XRD outcomes (Fig. 2), reaction (1) proceeded to a certain extent at 1000 °C,  until  the  consumption of the initial precursors was 1200 °C and 1 h of dwell range 1000-1200 °C,  completed at  time. Within  the  temperature  several  chemical  reactions  take place, giving rise  to stable  intermediate  solid  by-products.  In  particular,  the  emergence  of  a  mixture  of  hafnium silicides, most  probably Hf5Si4 contrast to other  and Hf5Si3, IVB group transition metals  was  ascertained.  In  like  titanium and zirco nium,  reliable  thermochemical data for  the  family of  hafnium silicides are unavailable. However, having as sumed highly plausible the chemical affinity of Hf with  respect  to Zr, and hence of  the present  studied system  with  an  Zr-Si-B4C composition was calculated as well  one,  a multiphase  equilibrium  (Fig. 3). Such an  exercise  basically  intended  discerning  the  predomi nance  of  the  chemical  reactions  evolving  during  the  studied synthesis. The conversion rates do not change  appreciably by  increasing  temperature. Thus,  from a  thermodynamic point of  view,  it  can be  argued that  reaction  (1)  proceeds  independently  of  the  set  temperature. A specific  discussion will  be  presented  later.  Raising the processing temperature of test up to 1650 °C,  the PLSHT-n  the  end-products of  the  synthesis  vary slightly in composition and relative percentages.  The XRD pattern  of  sample PLSHT-5  is  shown  in  Fig. 4.  3.2. The reactive hot-pressed composite  3.2.1. Microstructure  The processing conditions of the PLSHT-5 treatment  were adopted for the ceramic synthesis during the reac tive hot-pressing. In fact, the as-ground powder mixture  was  loaded straight  into the hot-press  equipment and  heat-treated, in accordance with the thermal programme  shown in Fig. 1. The heating rate from about 900 up to 1450 °C was kept 10 °C min\\x001  as  low as  in order  to  quench any emergence of spontaneous self-combustion,  whilst  the external pressure was stepwise applied only 1450 °C for  once  the  synthesis  at  60 h had  finished.  0.0 0  0.5  1.0  1.5  20.  2.5  0  0.1  0.2  0.2  0.3  0.4  0.4  0.5  0.6  0.6  0.7  0.8  0.8  Added Si/vmol  m  o  l  1  Added B 4C /mol  Zr   Zr2Si  ZrC   SiC   ZrB2  Zr5Si3  ZrSi   Fig. 3. Multiphase equilibrium composition calculated at 25 °C, and  2.2 mol  Zr  as  starting  amount.  Silicon  and  boron  carbide  are  progressively added up to 0.8 and 1 mol, respectively.  27  29  31  33  35  37  39  HfB2  m -H fO2  SiC  HfC   H fC   a)  b)  2θ/degree  C  n u o  t  s  (  a  .  u  )  Fig. 4. XRD pattern from sample PLSHT-5 (1450 °C for 1 h): normal  view (a), expanded full scale (b).  27  28  29  30  31  32 33 34 2θ/degree  35  36  37  38  39  C  n u o  t  s  (  a  .  u  )  S i  H f  Hf  Hf  *Hf   *Hf  *Hf   H fB2  H fC   HfC  -3)  H fO2  H fO2  S iC   B4C  HfH1 .5  (H f,S i)  (H f,S i)  (Hf,S i)   (Hf,S i)   (H f,S i)   (Hf,S i)  (H f,S i)   -2)  -1)  -0)  Fig. 2. XRD patterns of sample PLSHT-0 (as-ground), PLSHT-1 (1000 °C), PLSHT-2 (1100 °C) and PLSHT-3 (1200 °C).  *Hf:  ICDD  38-1478 but  longer lattice parameters.  1872  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879      \\x0c', 'The terminal stage was 1900 °C for 10 min, 50 MPa of  applied pressure.  In accordance with the  full  completion of  reaction  (1), the predicted bulk density 9.51 g cm\\x003. The measured is g cm\\x003,  of  the  composite  bulk  density was  9.43  0.992 of  the  theoretical  relative density. The  XRD analysis  of  the  hot-pressed  composite  (Fig.  5)  established the formation of HfB2, SiC, HfC, m-HfO2 as secondary phase, and a slight 0 0 l texture of the dib oride matrix. The estimated amounts (vol%) of the men tioned phases (Fig. 5) agree with those predicted on the  full completion of  reaction (1)  (72 HfB2, 22.1 SiC, 5.9 silicon introduced in the  HfC).  It  follows  that all  the  starting mixture reacts with the available carbon (from  B4C) and yields the expected amount of SiC. The lattice parameter of the cubic hafnium carbide, 0.46230 ±  0.0005 nm, slightly departs from the reference value of  0.463765 nm (ICDD 39-1491). Thus,  the occurrence of  a non-stoichiometric HfC is highly plausible.  The composite bulk,  inspected by SEM, does not re veal  residual porosity. A final  relative density of 0.992  congruently agrees with a microstructure free of poros ity. The  general  configuration  of  the microstructure  (Fig. 6(a)) presents regularly faceted diboride grains (from 0.5 to 4 lm as grain size) and intergranular SiC particulates (about 2 lm max as grain size). Considering  the relatively high temperature applied, the grain size re mains rather fine. The fracture mode is chiefly intergran ular even if,  in correspondence with the largest grains,  intragranular events  seem to prevail. Microcracking of  HfB2 grains and debonding of SiC/HfB2 observed. The emergence of tensile or compressive resid interfaces are  ual strained fields in the diboride matrix and SiC partic ulates, respectively, was considered the main reason of  such a phenomenon. Further SEM-EDX examinations  of polished regions identified the main solid constituents  (Fig.  6(b)). The  grain  size  of  SiC varies within  few  micrometers. The diboride/diboride interfaces seem de pleted of secondary phases.  Differently from previous  studies on HfB2-SiC sysincluding small quantity of additives [13,14], the  tems  reaction products associable to a residual  liquid phase  were not  found. This corroborates the concern that so lid-state diffusion drove  the  transfer of matter during  sintering, and that undesired thermally unstable second ary phases were most  likely suppressed.  3.2.2. Mechanical properties  The experimental values of some mechanical proper ties are summarized in Table 2. In terms of absolute vathe YoungÕs modulus (E) agrees with that reported  lue,  20  25  30  35  40  45  2θ/degree  C  n u o  t  (  a  .  u  )  Hf B2  mHfO 2  SiC   HfC  HfC  Hf B2  Hf B2  Fig.  5. XRD pattern  of  the  reactive  hot-pressed  composite:  the  volumetric amounts of HfB2 (71%), SiC (21%), HfC (7%), and HfO2  (1%) were  estimated,  using  a Rietveld routine. The  0 0 l  crystallo graphic orientation of the HfB2 matrix is evident.  Table 2  Mechanical properties of the reactive hot-pressed composite: YoungÕs  modulus E, micro-hardness Hv1.0, 4-pt. flexural strength r at 25 and 1500 °C  E (GPa)a  Hv1.0 (GPa)b  r (MPa)b  25 °C  1500 °C  520 ± 4  19.0 ± 1.0  770 ± 35  310 ± 15  a Uncertainty.  b  (mean ± 1 SD).  Fig. 6. SEs-SEM micrographs  from a fracture (a) and a polished surface (b) of  the reactive hot-pressed composite.  (a) The arrows mark the  imprinting of SiC particles upon HfB2 grains.  (b) Some grains of HfC (1) and HfO2  (2) are indicated; darker micrometric features consist of  intergranular SiC particles.  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  1873    \\x0c', 'for a similar fully dense HfB2-20 vol% SiC ceramic [3]. The flexural strength at room temperature exhibits a  really impressive mean modulus of rupture for this cat egory of UHTCs. Likewise,  the narrow dispersion of  data points out  the suitability of  the processing proce dures which enable 1500 °C,  some  control  over manufacture  flaws. At  a  reduction  in  strength  occurred.  The  load-displacement  curve deviates  to some  extent  from linearity (Fig. 7). The micro-hardness  (Hv1.0)  is  rather high:  the presence of hard ingredients like HfB2 not adversely affected either by residual  and  SiC is  porosity or by a coarse microstructure.  3.2.3. The resistance to oxidation  The mass change (w) vs. temperature (T) or vs. expo sure time (t) during the treatment T-1 is shown in Fig. 8.  The pattern of the thermogravimetric (TG) curve along  the heating stage (Fig. 8(a))  is closely connected to the  thermal  instability of  the oxide scale growing upon the  external  faces of  the oxidizing sample. The temporary 1400 °C of  stationary  branch  between  1280  and  the  TG data basically results from a balance between the re lease of gaseous by-products like B2O3 and CO, and the transformation of HfB2 and HfC into HfO2. An offset of 0.35 mg cm\\x002, which accounts for the oxidation preced ing the planned exposure, was subtracted from the TG  data  over  the  isothermal  stage  (Fig.  8(b)). Applying  the model proposed by Nickel  [15],  the multiple linear regression of  the paralinear  law w = KPAR  p  t + KLINt  (KPAR and KLIN constants) fits This calculation points out that  the TG data very well.  the oxidation process  is  rate-limited  by  diffusional mechanisms. The XRD  analysis  of  the  exposed  surfaces  detected monoclinic  HfO2, highly textured HfSiO4, and SiO2 cristobalite in minor content. The SEM-EDX inspection of the  cross-section  (Fig.  9(a))  highlights  an  external  scale,  basically consisting of a silica glass  (Fig. 9(b)). Local  ruptures of  the external oxide scale were not observed.  The undulating thickness of such a glassy coating indi cates that, owing to a diminished viscosity at the testing  temperature,  it may laterally flow out. Underneath this  glassy  layer,  an  oxide  scale  extending  up  to  the  as sintered virgin bulk is basically constituted of HfSiO4 and monoclinic HfO2. Both these oxides crystals are enclosed by a glassy melt, the former lying along the  bottom of the outermost silica glass.  As  far as  the oxidation treatment T-2 is concerned, (mg cm\\x002)  the  mass  gain  were  1.60 ± 0.05  and  1.85 ± 0.05  after  10  and  10 + 10 min  of  exposure,  respectively. The  scheme of  the microstructural  alter ation (Fig. 10) differs a little  from that  just described  (Fig. 9(a)). Apart  from the  external glass, only HfO2 crystals embedded in the glassy melt compose the under lying scale. Underneath this oxide scale, an additional  0  50  100  150  200  250  300  350  0  0.05  0.1  0.15  0.2  0.25  D/mm   M  O  R  /  M  P  a  Fig. 7. Load (MOR) vs. displacement (D) curve, recorded at 1500 °C  during 4-pt flexural strength test.  0 600  0.1  0.2  0.3  0.4  800  1000  1200  1400  T /˚C  w  /  m  g  c  m    2  w  /  m  g  c  m    2  (a)  0  0.2  0.4  0.6  0.8  1  1.2  1.4  1.6  0  200  400  600  800  1000  1200  t / min  Regression fit  Experimenta l data  KPAR= 0.17 mg2 cm -4 h-1 KLIN=0.0168 mg cm-2 h-1  (b)  Fig. 8. Thermogravimetric data of the treatment T-1: weight change w vs. temperature T (a) or vs. exposure time t (b). The best fitted parameters  KPAR and KLIN are shown.  1874  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879        \\x0c', 'porous band has  formed in consequence of  the active  oxidation of  the SiC particulates herein located. After  10 + 10 min of exposure, the thickness of specimen had an increase of about 20 lm.  the oxidized  4. Discussion  4.1. The ‘‘in situ’’ synthesis of  the HfB2-SiC-HfC  mixture  The ‘‘in situ’’ synthesis of the HfB2-SiC mixture proceeded via decomposition and further mass exchange  processes at high temperatures. Of course the solid pre cursors (i.e., Hf, Si and B4C) were selected such that the resulting reaction was highly exothermic. The real driv ing force of such solid-state exchange reactions definitely  arises  from the  formation of  thermodynamically very  stable end-products.  The  strong  exothermicity of  the  studied system of  reactants  switched a sort of mechanically induced for mation of by-products during ball-milling. In particular,  apart  from a predictable reduction in size of  the ball milled powders, an accumulation of defects in the pow der particles introduces additional energy to the reactant  system in the  form of  interfacial  and strain energies.  Such extra available energy leads part of the intimately  mixed  components  to  chemically  react  in  this way,  2Hf + B4C = 2HfB2 + C, without thermal inputs. The formation of an inert  the help of  external  compound  like HfB2 plausibly quenches other reactions from selfpropagating during mechanical mixing. Likewise, an 10 °C min\\x001  heating  rate  of  during  PLSHT  tests  smooths any occurrence of spontaneous combustion.  The multi-component equilibrium composition (Fig.  3)  along with  the  support  of XRD analyses  on  the  PLSHT samples, helped in tracing a sequence of some  transitory steps of reaction (1). By analogy with a simi lar system [16], the following chemical reactions (dG25: Gibbs free energy of reaction at 25 °C) 2Hf þ B4C ¼ HfB2 þ HfC ðdG25 ¼ \\x00794 kJ=molÞ  ð2Þ  xHf þ ySi ¼ Hf xSiy ;  x=y ¼ 5=4; 5=3  ð3Þ  2HfC þ 3Si þ B4C ¼ 2HfB2 þ 3SiC ðdG25 ¼ \\x00386 kJ=molÞ  ð4Þ  HfC þ Hf 5Si4 þ 3B4C ¼ 6HfB2 þ 4SiC  ð5Þ  are guessed to occur. A lack of thermochemical data for  the Hf-Si  system does  not  permit  the  dG values  of  reaction (3) and (5)  to be calculated, even if,  similarly  to Zr-Si or Ti-Si systems,  large negative dG values are  expected. Likewise,  the predominance of  specific haf nium silicides at different temperatures logically depends  on the activity ratio aHf/aSi of the available metallic Hf and Si. Also, reaction (4) and (5) describe the delayed  formation of SiC, only once an appreciable amount of  HfC has already formed. The XRD pattern of  sample  Fig. 9. SEs-SEM micrograph from a polished cross-section of the hot pressed composite, after oxidation at 1450 °C for 20 h (a), and EDX spectrum  ((b) 4 keV electron beam energy) of the external glass (upper part of the micrograph).  Fig. 10. SEs-SEM micrograph from a polished cross-section of the hot 1700 °C for 2 · 10 min. The  pressed composite,  after oxidation at  upper part of  the micrograph is occupied by  a  silica based glassy  coating.  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  1875  \\x0c', 'PLSHT-2  show robust  peak  intensities  belonging  to  HfC and some hafnium silicides (Fig. 2). Increasing temperature up to 1200 °C, reaction (4) and (5) become pre vailing, giving rise to the desired formation of HfB2 and SiC.  This  representation intended to highlight  the domi nant chemical reactions. The finite size of the powdered  precursors, for instance, often put in contact only some  of  them,  favouring reactions which may yield minority  products. The XRD patterns of  sample PLSHT-2 and  PLSHT-3  still  contain minor  unindexed  peaks. The  application of synthesis 1650 °C does not  temperatures  from 1300 up to  imply significant changes  in the stoi chiometry of  the end-products  (Table 1). Nevertheless,  the reduction of  the peak width reflects the predictable  increase in size of HfB2, SiC and HfC crystallites. Moreover, the gradual disappearance of the HfO2 phase is explained by its conversion into HfC. The carbon-rich  reducing  environment  favours  the  transformation  of  HfO2 into HfC, according to the following reaction HfO2 þ 3C ¼ HfC þ 2COðgÞ  ð6Þ  which has negative dG above 1375 °C. Finally, for temperatures above 1600 °C, the conversion yield of HfC starts dominating over that of SiC. Therefore, 1450 °C  was deemed an acceptable temperature for synthesizing  as fine as possible solid constituents  for  the final com posite and, at  the same time,  for promoting the forma tion of HfB2 and SiC against HfC.  the  competing raise of  4.2. Reactive hot-pressing, microstructure and mechanical  properties  Compared to the hot-pressing conditions of an addi(2200 °C  tive-free HfB2-20 vol% SiC powder mixture for 1 h and 25 MPa [3]), the present material  experi enced a less  intensive  thermal history. This  represents  a very important issue because the intrinsic poor sinter ability of a mechanically mixed HfB2-SiC powder system is overturned by a correct selection either of the  precursors and of the heat  treatments.  The initial fineness of  the powder mixture is decisive  in obtaining well dispersed and fine-grained solid phases  constituting the sintered composite. Once B4C decomposes, boron and carbon readily diffuse and interact  with the metallic precursors Hf and Si. This  feature is  supported by the evidence that the ‘‘in situ’’ synthesized  solid HfB2 and SiC keep the dimensional ranges of starting precursors.  the  The densification rate of diboride powders is known  being greatly influenced by the boron activity. A con tamination in oxygen,  typically present upon the dibo ride particle  surfaces,  entails a decrease  in the boron  activity (i.e.,  reduction of  the densification rate), and  brings on grain-coarsening at  the processing tempera ture required for achieving near full density [17,18]. Dif ferently  from  HfB2 manipulated in air before hot-pressing  powders  usually  stored  and  [3,5,13,14],  in  the present  study the surfaces of  the HfB2 particulates synthesized in a inert atmosphere) are sup (‘‘in situ’’  posed being only slightly affected by the oxygen contam ination.  Further  conditions  of  an  intimate  contact  between oxygen-depleted reacting interfaces of diborides  and the contemporary application of high temperature  and pressure most likely facilitate surface/boundary dif fusion, and assist densification, the boron activity being  no further depressed.  In addition, similarly to an analogous ZrB2-SiC system [19], the acceleration in densification is compensated  by the dragging action of SiC, which, by virtue of  its  own refractoriness,  limits extensive growth of  the HfB2 imprinted all around the HfB2 convincing evidence of the role of SiC  matrix. Some grooves  grains may be  in controlling growth of  the diboride matrix (Fig. 6). 1900 °C,  Despite  a  sintering  temperature of  the final  average grain size of the tested material does not exceed  a few micrometers.  Clean diboride/diboride or diboride/carbide bound aries,  already  reported  for  an  ‘‘in  situ’’  synthesized  TiB2-SiC system [16], interfacial joinings. Such a merit  support  the supposition of  tight  can be  ascribed to  the high surface  energy of  the newly formed diboride  grains, and indicates  that  the  tightening SiC particles  formed in situ strongly connect with the diboride matrix.  This inference should be validated by assessing at nano metric level  the absence of  foreign phases along grain  boundaries,  that  analyses  (herein  presented)  verified  only with the SEM-EDX technique.  As far as the mechanical characterization is concerned, a YoungÕs modulus of 520 GPa is in very good  agreement with that  calculated (i.e., 505 GPa) using a  rule  of mixture  [20],  namely  the  arithmetic  average  RQiEi, Qi  the volume  fractions of phases  constituting  the composite. For Ei and Qi used: 530 GPa [21] and 0.71  the following data were  for HfB2, and 0.06 for HfC, 448 GPa [23] and 0.22 for SiC, and  460 GPa  [22]  239 GPa [21] and 0.01 for HfO2. Control of the residual porosity along with of the secondary phases character ized  by  low E values  (i.e., HfO2) importance against the depression of  is of  fundamental  this property.  In the results section, the opening of HfB2/SiC interfaces and some microcracking of HfB2 grains were reported occurring after cooling. A representative  example is shown in Fig. 11. The extent of  the residual  thermal  stresses which involve the HfB2/SiC interfaces was calculated in accordance with the formula proposed  by Eshelby [24]. It basically considers spherical particles  embedded in an infinite matrix (in our  case SiC and  HfB2, volume fractions of the phases constituting the compos respectively) but does not  take into account  the  ite.  Actually,  as  far  as  the  HfB2-SiC  system  is  1876  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  \\x0c', 'concerned,  referenced assignments of  the freezing tem perature value TFR, which commonly indicates the temperature range along the cooling stage within which the  stresses  are not  released by atomic diffusion, are not  available. Differently from other to 1200 °C [6],  authors which  set  TFR equal an HfB2-SiC system is plausible to call for the input of larger TFR value around 1800 °C. Magley et al. [25] utilized a TFR value of 1700 °C for a SiC-TiB2 particulate composite. Another proposal modelled by Mizutani  the  strong refractoriness of  on a SiC-rich ZrB2 composite extrapolated a value of TFR equal to 1750 °C [26]. On substitution of the appropriate values, tensile stressed HfB2 and compressive expected (Table 3). Therefore,  stressed  SiC are  the  microcracking events  are attributed to the  emergence  of high residual  stresses which come up at  the HfB2/ SiC interfaces in consequence of the mismatch of structural properties like E, thermal expansion (k) and PoissonÕs ratio (m). Such a first approximation, based on the  HfB2 matrix - SiC particle model, overlooks ence of HfC. On the other hand, only HfB2 and SiC were involved in these localized micro-mechanical phe the pres nomena. In addition, (TFR = 1800 °C, 10\\x006/°C\\x001 [22])  the input of  the following values m = 0.18, k = 6.5 ·  E = 460 GPa,  in the Eshelby formula foresees  com pressive radial stresses for HfC (\\x001.095 GPa) and tensile stresses (0.55 GPa) for HfB2. It follows that, in the economy of this basic estimate, the presence of 6 v/o HfC can  realistically be overlooked.  The flexural strength and relative dispersion (Table 2)  exhibit excellent values, 1500 °C. At 1500 °C,  at  room temperature  and  the  refractoriness of HfB2 composite skeleton, ensures  and  SiC, which compose  the  the strength is  retained efficiently. Without any plastic  deformation  of  the  test  bar  taking  place,  the  slight  downward  curvature  of  the  load-displacement  curve  (Fig. 7)  is  connected to the sub-critical crack growth. the surfaces fractured at 1500 °C  SEM observations of  confirmed the occurrence of this phenomenon. Actually,  the microstructural degradation on the specimen tested 1500 °C adversely  at  contributes  to  the  decrease  in  strength. Most of the HfC grains directly facing the oxi dizing atmosphere for instance react readily, transform ing into an oxide form. However, the inner parts of the  specimen are protected from further attack by the ability  of  the outermost oxidation product  like a silica-based  glass  to seal  the external  surfaces,  thus preventing the  strength being further impaired.  The  initial  compositional  design  of  the  composite  aimed  at  introducing  a moderate  level  of HfC in  a  HfB2 + SiC skeleton. This specific feature was reported to have succeeded in improving oxidation/ablation resis tance of diboride-SiC materials subjected to heating re gimes  simulating  hypersonic  re-entry  space missions  [27]. However, with reference to the strength behaviour in air up to 1500 °C (conditions  that differ markedly  from those  just mentioned),  increasing the  content of  HfC negatively affects the ability of HfB2-SiC composites to retain the original strength effectively.  4.3. The resistance to oxidation  The present study highlighted how limited mass gains  are accompanied by a few alterations of  the original  microstructure.  In  agreement  with  other  authors  [10,11,27,28], the presence of SiC particles substantially enhances the resistance to oxidation of a pure HfB2. ItÕs known that the oxidation of HfB2 and HfC generates HfO2 and amorphous B2O3, or HfO2 and CO(g), respectively. The increase in the scaling rate above 800 °C (Fig.  8) is primarily caused by the selective oxidation either of  the HfB2 and HfC, with no appreciable attack of the SiC particles. The external oxide scale that forms partially  allows  the diffusion of oxygen through interconnected  pores  or  via  lattice  vacancies  in HfO2. ZrO2, HfO2 behaves plausibly as an anionic conductor. In addition, at a relatively high temperature, a fluid B2O3 (melting point of 450 °C), which most likely covers the external faces of the oxidizing sample, is known to  Similarly  to  be much more permeable to oxygen than a silica glass  [29].  The  almost  stationary  trend  of  the  TG data  Fig.  11. SEs-SEM micrograph from a polished section of  the  as sintered  composite:  black  and  white  arrows mark  examples  of  disjoined  HfB2/SiC  interfaces  and  microcracked  HfB2  grains,  respectively.  Table 3  Estimation  of  the  residual  stresses  (rRES)  at  the matrix/particle  interface (E, YoungÕs modulus; k,  linear thermal expansion coefficient;  m, PoissonÕs ratio)  E (GPa)  k 10\\x006 (°C)  m  rRES (GPa)  HfB2 (matrix)  530  8  0.12 [21]  1.407 \\x002.815  SiC (particle)  448  4  0.168 [23]  F. Monteverde / Composites Science and Technology 65 (2005) 1869-1879  1877  \\x0c', 'between 1280 and 1400 °C (Fig. 8) accounts for the bal anced competition between mass losses and gains.  Substantial benefits from the presence of the SiC partemperatures above 1400 °C.  ticles  result  for  In accor dance  with  the  following  reaction,  SiC + 3/2  O2(g) = SiO2 + CO(g), silica which forms from the oxidized SiC particles combines with the available boria,  providing more  oxidation  protection  than  the HfO2 alone. Likewise other MB2-SiC systems, M = Zr or Hf [10,11,28,30,31], the exposed faces of the present com posite  are  covered  by  an  adherent  silica-based  glass  (Fig. 9(b)), which in turn is characterized by an undulat ing thickness. This glass, scarcely permeable to oxygen  at  the  tested temperature  [29],  seals  and protects  the  external surfaces of the sample. The evidence of residual  unoxidized SiC particles just beneath the external oxide  scale  settles  the  fundamental  role of  such a phase  in  slowing the advance of the oxidation attack (Fig. 12).  TG data of the isothermal test T-1 fit a paralinear law  very well (Fig. 4). The parabolic contribution dominates  the monotonically decelerating trend of the TG data. In  effect,  the growth of a protective  external oxide  scale  progressively imposes longer diffusion paths for oxygen  to arrive at  the diboride-oxide interface. The departure  from a pure parabolic pattern is described by a little neg ative linear contribution, and is most likely motivated by  a release of  some  gaseous oxidation by-products  like  CO(g) and B2O3(g). Moreover, of the diboride-diboride boundaries  the apparent cleanness  (i.e.,  absence  of  intergranular compounds) has a beneficial merit in lim iting  the  preferential  inward  transport  of  oxygen  through them (Fig. 12).  As far as the treatment T2 is concerned, the intrinsic  refractoriness of  the studied system and the protection  of  the external glassy coating supply the thermostruc tural  stability that enables  the composite to withstand  efficiently such severe thermal  loads. In addition, the in crease of about 20 lm in thickness of the oxidized sam ple after 10 + 10 min of  exposure verifies  that  surface  conversion and removal of mass have no appreciable  effects.  5. Summary  This work highlighted promising advances in the ‘‘in  situ’’  synthesis, microstructure and mechanical proper ties of an ultra-high-temperature HfB2-SiC composite. Solid powder precursors like Hf, Si and B4C were properly processed via reactive hot-pressing, and a full dense  HfB2-SiC composite thus obtained. The full conversion of the starting reagents into the end-products (HfB2, SiC, HfC) was ‘‘in situ’’ conducted during the reactive run at 1450 °C for 60 min. The final microhot-press structure, mean grain size 3 lm, was uniform and rather  fine, with HfO2 as SiC phase basically played the role of  the principal  secondary phase. The  inhibitor against  excessive  coarsening of  the HfB2 matrix. The flexural very promising: 770 ± 35 and 1500 °C, 310 ± 15 MPa at 25 and respectively. The YoungÕs modulus was 520 GPa. The composite showed  strength  values were  a rather good resistance to oxidation: repeated exposures at 1700 °C in air did not severely affect the overall  integrity  of  the  composite. This  unconventional  ap proach  of  ‘‘in  situ’’  synthesizing/densifying  strongly  covalent  ceramic-matrix composites was an novel an swer to fabricating ultra-refractory materials.  Acknowledgements  I  sincerely acknowledge  the helpful  contribution of  the colleagues D. Dalle Fabbriche (thermal treatments),  C. Melandri (mechanical tests), and A. Balbo (oxidation  tests).  References  [1] Upadhya K, Yang J-M, Hoffman WP. Materials  for ultra-high  temperatures  structural  applications.  Am  Ceram  Soc  Bull  1997;58:51-6.  [2] Wang CR, Yang  J-M, Hoffman WP.  Thermal  stability  of  refractory  carbide/boride  composites.  Mater  Chem  Phys  2002;74:272-81.  [3] Gasch M, Ellerby D, Irby E, Beckman S, Gusman M, Johnson S.  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Monteverde / Composites Science and Technology 65 (2005) 1869-1879  1879  \\x0c']"
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  "_id": 226,
  "PDF": "Promising UltraHighTemperature Ceramic Materials.pdf",
  "Text": "['ISSN 0036(cid:2)0236, Russian Journal of Inorganic Chemistry, 2013, Vol. 58, No. 14, pp. 1669-1693. © Pleiades Publishing, Ltd., 2013.  Promising Ultra(cid:2)High(cid:2)Temperature Ceramic Materials  for Aerospace Applications  E. P. Simonenkoa, b, D. V. Sevast’yanova, N. P. Simonenkoa,  V. G. Sevast’yanova, and N. T. Kuznetsova  aKurnakov Institute of General and Inorganic Chemistry, Russian Academy of Sciences,  Leninskii pr. 31, Moscow, 119991 Russia bLomonosov State University of Fine Chemical Technologies, pr. Vernadskogo 86, Moscow, 119571 Russia  Abstract—Some aspects of heat transfer upon the interaction between components with a sharp leading edge and high(cid:2)enthalpy high(cid:2)speed flow of dissociated air have been considered; some material characteristics, which should be primarily taken into account when prognosticating the behavior of materials that are prom(cid:2) ising for using as components of hypersonic flight vehicles, have been substantiated; specific features of the oxidation of materials based on zirconium and hafnium diborides have been touched briefly; the methods of increasing oxidation resistance of these materials that have been developed by various groups of researchers have been demonstrate; some works concerning the behavior of samples under the effect of high(cid:2)enthalpy flows of dissociated air have been described, including those that simulate sharp leading domes and edges of wings of hypersonic flight vehicles.  DOI: 10.1134/S0036023613140039  The solution of one of the most topical problems of modern structural(cid:2)materials science, i.e., the develop(cid:2) ment of the technologies of materials capable of oper(cid:2) ating at ultra(cid:2)high temperatures under the conditions of aerodynamic heating, which are important for the realization of a breakthrough in the creation of hyper(cid:2) sonic flight vehicles and in rocket engineering, is con(cid:2) nected with the fundamental problem of the direc(cid:2) tional selection of compositions and development of methods of producing ultra(cid:2)high(cid:2)temperature materi(cid:2) als. For decreasing aerodynamic drag and increasing lifting force and maneuverability, the supersonic flight vehicles must have profiles with sharp leading edges (radius of curvature from a few tenths of a millimeter to several millimeters). As a result, the article is sub(cid:2) jected to the action of powerful heat fluxes at the stag(cid:2) nation point (several MW/m2); correspondingly, the temperature of the surface can become more than 2000°C [1-20] (up to 2500-2600°C), which exceeds the exploitation temperature of standard materials, such as composites with a silicon carbide (SiC) matrix reinforced by carbon fibers (C/SiC). At so high tem(cid:2) peratures, the protective layer of silicon carbide on C/SiC composites becomes chemically active [21- 25], which determines the rapid ablation of the mate(cid:2) rial and the leveling(cid:2)off of the protective properties of SiC with respect to the reinforcing carbon fibers. The use of materials with a high thermal conductivity ensures a more rapid redistribution of the excessive heat and favors its removal from the line of the total stagnation of the air stream in the case of components with sharp leading edges.  Thus, for substances could be applied as compo(cid:2) nents of ultra(cid:2)high(cid:2)temperature ceramics (UHTCs), they must possess not only high melting point, but also high phase stability in a wide temperature range (from room temperature to the expected service tempera(cid:2) ture),  good mechanical characteristics  (including those of the products of their oxidation), and suffi(cid:2) ciently high thermal conductivity (see Table 1). In the literature, the primary attention is given to refractory borides of transition metals (ZrB2, HfB2, which have the greatest oxidation resistance in the row of borides of transition elements) with additions of refractory carbides (SiC [2-12, 14-20, 26-107], TiC, ZrC, HfC, TaC [1, 3-5, 12, 44, 62, 79, 81, 86, 89, 92, 98, 104, 108-114]), silicides (MoSi2, TaSi2, ZrSi2 [1, 2, 15, 17, 32, 48, 64, 89, 91, 95, 100, 111, 115-118]), borides (TaB2, CrB2, LaB2, WB2 [2, 10, 11, 19, 51, 57, 76, 80, 112, 109, 119]), nitrides [40, 70, 71, 84], and refractory metals (in particular, Ir [2, 32, 120]), which is connected with the attempts to increase the oxida(cid:2) tion resistance of materials based on the borides of zir(cid:2) conium and hafnium.  The large contribution of Russian (both Soviet and post(cid:2)Soviet)  researchers  [121-122], as well as of European and American scientists [123],  into the development  of  ultra(cid:2)high(cid:2)temperature materials intended for the extreme operating conditions should be noted; in the recent decades, Chinese researchers begin also to play larger and larger role.  The bibliographical search with the use of the Sci(cid:2) Finder (CAS) system based on keywords that charac(cid:2) terize this field of research made it possible to obtain a  1669  \\x0c', '1670  SIMONENKO et al.  Table 1. Some properties of refractory compounds that are promising as components of UHT ceramic composites  Compound  Melting  temperature, °C  Density, g/cm3  Thermal conductivity,  W/(m K)   Thermal expansion coefficient, 1/K  HfC  3890 [124]  12.65 [125]  20 (323 K) [127]  TaC  ZrC  HfB2  3880 [124]  3983 [128]  3427 [130]  3500 [131]  3530 [124]  3250 [134]  14.64 [129]  14.58 [125]  6.54 [132]  6.60 [133]  11.18 [135]  11.21 [136]  11.20 [137]  22.2 (300 K) [124]  20.5 (300 K) [124,  131]  51.0 (300 K) [138]  60.0 (1300 K) [138]  143.0 (2300 K) [138]  ZrB2  3040 [134]  3000 [140]  6.09 [141]  6.00 [142]  58.0 (300 K) [138]  64.5 (1300 K) [138]  134.0 (2300 K) [138]  6.19 × 10-6 (300-1273 K) [126] 7.10 × 10-6 (300-2273 K) [126] 7.54 × 10-6 (300-2873 K) [126]  6.61 × 10-6 (300-1273 K) [126] 7.31 × 10-6 (300-2273 K) [126] 7.80 × 10-6 (300-2873 K) [126]  6.73 × 10-6 (300-1373 K) [122]  6.3 × 10-6 (300-1300 K) [138]; 6.8 × 10-6 (1300-2000 K) [138]; 6.64 × 10-6 (1100-1800 K, [100]) [139]; 6.89 × 10-6 (1100-1800 K, [001]) [139]; 6.39 × 10-6 (300-1000 K, [100]) [139]; 6.81 × 10-6 (300-1000 K, [001]) [139]  5.9 × 10-6 (300-1300 K) [138] 6.5 × 10-6 (1300-2000 K) [138] 5.92 × 10-6 (1100-1800 K, [100]) [139] 8.62 × 10-6 (1000-2300 K, [100]) [139] 7.65 × 10-6 (1000-2300 K, [001]) [139]  MoSi2  TaSi2  ZrSi2  Ir  2030 [143]  2088 [144]  2000 [145]  2399 [149]  2400 [124]  1620 dec. [151]  2450 [153]  6.27 rt [146]  6.28 rt [147]  6.41 ht [148]  9.07 [146, 150]  4.88 [152]  21.52 [154]  22.57 [155]  22.56 [156]  32.7 (922 K) [149]  14.4 (1478 K) [149]  8.2 × 10-6 (300 K, [100]) [145] 9.4 × 10-6 (300 K, [100]) [145]  -  -  147.5 (289.2 K) [157]  138.5 (492.2 K) [157]  139 (500 K) [157]  8.5 × 10-6 (300-1750 K) [149]  -  9 × 10-6 [153]  SiC  2824 [158]  3.2 [159, 160]  -  (4.7-5.5) × 10-6 (300-1773°C) [159] 5.12 × 10-6 (300-1300 K) [161]  distribution of the frequency of references to different chemical compounds (Fig. 1). As can be seen, the most frequently mentioned compounds are ZrB2 and HfB2, as well as SiC, which are the basic components of  ultra(cid:2)high(cid:2)temperature  ceramics  (UHTCs). Although  the melting  points  of  zirconium  and hafnium carbides are substantially higher than those of the corresponding diborides, the latter are considered as more promising for the above(cid:2)mentioned applica(cid:2) tions because of their higher thermal conductivity, especially, at high temperatures (see Table 1). Further(cid:2) more, the materials based on ZrB2 and HfB2 have very good mechanical characteristics: high hardness [162, 163], high strength (at room temperature, values of 500 MPa and above have been observed [14, 17, 26, 65, 66, 68-70, 72, 73, 80, 81, 83, 117, 118]); the modulus of elasticity at room temperature is ~400-500 GPa [14, 26, 68, 69, 76, 80, 117].  Of large importance is also the fact that upon the oxidation of ZrB2 and HfB2 there arise, apart from products that are liquid at temperatures higher than 800-1200°C  (forming  borosilicate  glass, which impedes rapid diffusion of oxygen into the bulk of the material), also refractory oxides (Tm ~ 2700°C for ZrO2 [164, 165], 2800-2850°C for HfO2 [166, 167]) possessing a low vapor pressure (MO(g) at T = 2160°C is ~2.88 × 10-7 atm for ZrO2; ~1.15 × 10-7 atm for HfO2 [168]) and a sufficiently high (for the expected operating conditions) mechanical strength; this must ensure a maximally possible retention of the geometry of an article, which influences its aerodynamic char(cid:2) acteristics.  In the field of designing UHT materials, two direc(cid:2) tions, which are substantially different in, first of all, the priority purposes, can be distinguished. The first of  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1671  Number of publications for the compound, %  0  5  10  15  20  25  30  35  SiC ZrB2 HfB2 ZrO2 C ZrC Graphite SiO2 Al2O3 Si B4C Zr HfC MoSi2 TaC AlN Hf B2O3 Y2O3 Si3N4 HfO2 B BN Mo Ta Ni  Fig. 1. Distribution (%) of the publications concerning the ultrahigh(cid:2)temperature materials indexed over the com(cid:2) pounds indicated (according to the data of SciFinder, CAS) .  these is connected with the modification of classical UHT materials (Cf/C and Cf/SiC composites) by refractory compounds,  such as ceramic matrices and/or protective (antioxidant) coatings, with the pur(cid:2) pose to significantly increase their oxidation resistance without a significant loss of the excellent thermome(cid:2) chanical characteristics of the base materials [3, 4, 51, 57-64, 88, 108-110, 116, 119, 170-173].  An alternative direction is the creation of purely ceramic almost poreless materials (possibly, with a certain  their strengthening by  the  introduction of nanostructured components [33, 90, 97, 102, 103, 174, 175] or of short fibers [50, 65, 66, 70-73, 96, 97]) on the basis of substances with a high thermal conduc(cid:2) tivity. This will allow, as was already said above, an effi(cid:2) cient heat removal from the regions adjacent to the  points or lines of the total flow stagnation, which will work as separate segments of a complex composite detail (Fig. 2) [11, 20], whose different parts will bear different types of load—thermal, oxidation, shock, shear, etc. An additional argument in favor of the selection of the design of the thermo(cid:2)loaded parts of the flight vehicles in the form of an assembly of several segments is that the UHTCs are brittle ceramic mate(cid:2) rials  (the probability of  their destruction can be described in terms of the Weibull statistics); therefore, from the viewpoint of repairability, structures are pref(cid:2) erable that have a limited size or volume of compo(cid:2) nents and, therefore, have a reduced probability of the appearance of defects, which  can decrease  the strength. The components of small size can be pre(cid:2) pared more easily, especially with the use of methods of hot pressing or sintering in an electric field. If the size of separate parts is several centimeters, then the components such as leading edges of the wings or nose domes must consist of  segments. The  interfaces between the segments and the sites of junction with other components are very important from the view(cid:2) point of construction [177]. Such segmental structures have been developed within the NASA program of the “Next generation launch technology” (NGLT) [20]; examples of segmental components of the  leading edge of a wing are given in Figs. 2 and 3.  In  spite  of  the  conspicuous  difference  of approaches, the basic roles in them, on the whole, are played by the same inorganic compounds, namely, the super(cid:2)refractory borides and carbides of zirconium, hafnium, tantalum and some silicides; therefore, the principal attention  in  this communication will be given to the topical problems of the creation of ultra(cid:2) high(cid:2)temperature ceramic materials (UHTCs).  In this paper, we consider some aspects of heat transfer under the conditions of interaction between the components with a sharp leading edges and the high(cid:2)enthalpy high(cid:2)speed  flows of dissociated air (which makes it possible to substantiate a number of properties of the materials that should primarily be considered upon the prognostication of the behavior of materials that are promising as the components of  (a)  Metallic leeward  skin  Metallic structural elements  Thermal mass and/or radiation shield  Ti Structure and skins  (b)  ZrO2 Insulator  C/SiC Structure  HfB2/SiC Segmented leading edge  Hot structure attachment Leeward  UHTC segmented leading edge components  6 in.  Si3N4 Thermal mass  CMC TPS  15.5 in.  7.1 in.  C/SiC Structure  Windward  Carbon composite windward skin/TPS  Fig. 2. Construction of the leading edge of a wing prepared from a UHTC [11, 20].  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1672  SIMONENKO et al.  Fig. 3. Segmented components of the leading edge of a wing prepared from a UHTC [20].  Flow around shock wave  Free stream flow  (a)  qrad  qrad  qconv  qchem  (b)  Surface material  Insulation  Radiation back to the environment  Attachment region remains relatively cool  High heat flux to the stagnation region  Conduction away from the nose  Fig. 4. A schematic of the processes of aerodynamic heating and redistribution of the heat in hypersonic flight vehicles with (a) “blunt” [39] and (b) “sharp” [13] leading edges.  hypersonic flight vehicles); touch specific features of the oxidation of materials based on the zirconium and hafnium diborides; demonstrate methods of increas(cid:2) ing the oxidation resistance of these materials devel(cid:2) oped by different groups of researchers; and describe some works concerning the behavior of samples under the action of high(cid:2)enthalpy flows of dissociated air, including samples with a shape that simulates sharp leading parts and edges of wings of hypersonic flight vehicles.  1. ON THE HEAT BALANCE UPON  THE INTERACTION BETWEEN  COMPONENTS WITH A SHARP EDGE  AND HYPERSONIC FLOW [13, 17, 39]  Most clearly and briefly, the process of a redistribu(cid:2) tion of heat upon the aerodynamic heating of compo(cid:2) nents with sharp edges is described in [13, 17, 34, 39]. First of all, it should be noted that the term “sharp leading edge” refers to components, whose radius of curvature is substantially less than the basic overall size: for the so(cid:2)called “blunted” edges of reusable space Shuttle apparatuses there are characteristic radii of curvature of an order of tens of centimeters, while for the modern and promising concepts of hypersonic flight vehicles there are expected radii of curvature from millimeters to fractions of a millimeter. The  blunting of the edges was applied to soften the thermal load when overcoming the atmosphere upon re(cid:2)entry. Figure 4a illustrates the case of the removal of part of heat due to the arising shock wave; as a result, with taking into account the trajectory of motion, the max(cid:2) imum temperature of the nose parts and leading edges of the apparatus is limited to ~1650°C, however, this construction limits both the maneuverability, and the speed.  For the hypersonic flight vehicles, an improvement in the maneuverability requires that a laminar flow should exist over the entire surface; to this end, the application of sharp edges is necessary, which assumes the realization, first of all, of a convective heating of local regions to temperatures of ≥2000°C (Fig. 5 dis(cid:2) plays the result of the simulation of the heating of a component with a sharp edge by a high(cid:2)enthalpy flow of dissociated air (H0 = 16 MJ/kg) with the allowance for the partial oxidation of the initial material ZrB2- SiC [17]).  As a result, the heat supplied from the heated flow of gas (qconv) and from the thermal effects due to chemical reactions that occur on the surface (oxida(cid:2) tion of materials and the realization of processes of the recombination of the oxygen and nitrogen atoms cat(cid:2) alyzed by the surface (qchem) can be spent during the heat transfer from the most heated parts to less heated parts, where a reradiation can occur (qrad), and via its  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1673  nose section and wing edges there is formed a shock wave; when passing through this region, the gases are heated sharply and partially undergo dissociation and ionization, i.e., the molecules of O2 and N2 decom(cid:2) pose into the atomic oxygen (this process occurs pre(cid:2) dominantly because of the weaker bonds in O2 as com(cid:2) pared to those in N2) and nitrogen. Due to their exo(cid:2) thermic  nature,  the  reactions  of  the  catalytic recombination on the surface can make a significant contribution to the heating of the surface in addition to the basic contribution from convective heating at the stagnation point. The radiative heating from hot ionized gases in the flow often also plays a certain role in an increase in the temperature of the most thermo(cid:2) loaded local region. The convective component of heat absorbed by the surface (qconv) depends on the velocity, pressure, and temperature of the incident flow, as well as on the composition of the gas, geometry of the body, and on whether the flow is laminar or turbulent. The creation of turbulence as a result of the high surface roughness can considerably increase convective heating. In the incident flow, there is a relatively high con(cid:2) tent of dissociated gases (in Fig. 7, the relative con(cid:2) tents of O, O2, N, NO, and N2 at the temperatures of the surface of ~1800-2000 (L) and 2300-2600°C (H) are shown [10]), which not only can participate in the chemical reactions with the material (upon its oxida(cid:2) tion) but also can undergo  reactions of catalytic recombination with the formation of O2, N2, and NO molecules. In [48], it was shown that the atomic oxy(cid:2) gen is substantially more active chemically; the thick(cid:2) ness of the oxidized layer at the equal oxygen contents in the system is much higher in the case of O than in the case of O2 (according to [52], by a factor of 1.5 at the temperature of 1500°C and equal partial pres(cid:2) sures). All above(cid:2)mentioned processes are exother(cid:2) mic; as a result, additional energy can be transmitted to the surface. In the first approximation, the corre(cid:2) sponding thermal flow qchem is proportional to a coef(cid:2) ficient γ (which can vary from 0 to 1) that describes the total catalytic activity in the reactions of recombina(cid:2)  h  2R*  2450°C  2300  2200  2100  2000  1900  1800  1720°C  Fig. 5. Temperature distribution over the surface of a ZrB2-SiC sample with allowance for its oxidation accord(cid:2) ing to the data of simulation [17] (R* = 0.14 mm, h = 0.5 mm, and H0 = 16 MJ/kg).  transfer into the deeper parts of the material due to the thermal conductivity for putting(cid:2)on the active cooling system [176].  In the case of a sufficiently high thermal conductiv(cid:2) ity of the material, no failure from the thermal shock will occur in local regions upon a sharp increase in the temperature; the ceramic composite can pass the heat through itself, thereby removing it from the system (Fig. 4).  Squire and Marschall [13] described in detail some aspects of the interaction of hypersonic flows (with a Mach number >5) with the sharp edges of wings or with nose parts  in  terms of aerothermodynamics. Thus, in Fig. 6 it is shown that directly before the sharp  (a)  Sonic BL region  Sonic line Subsonic BL region  Boundary layer edge  Shock stand off  Free stream  Bow shock  Nose radius  Shock layer  (b)  T∞  Radiation  Boundary layer  Te  TW  Wind leading edge or nose tip  Free stream  Convection  Chemical heating  Conduction  Surface depth, x  Fig. 6. (a) Specific features of a hypersonic flow and (b) the energy balance on the surface [13].  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1674  SIMONENKO et al.  n  o  i  t  c a  r  f  e  l  o  M  1.0  0.8  0.6  0.4  0.2  0  L  H  L  H  L  H  L  H  L  H  L H  N  N2  NO  O  O2  1.5  1.6  1.8 3.3 Test specimens  3.5  3.9  Fig. 7. Molar fractions of N, N2, NO, O, and O2 arising in the near(cid:2)surface layer of the sample in the case of compar(cid:2) atively  low(cid:2)enthalpy (L) and high(cid:2)enthalpy (H) treat(cid:2) ments by gas flows [10]. As can be seen, virtually no molec(cid:2) ular oxygen remains when using both regimes of the tests, while the degree of dissociation of N2 grows sharply in the case of the high(cid:2)enthalpy regime.  tion. Furthermore, upon the quantitative simulation of the processes, the concentrations and the molar masses of the reagents, as well as the dissociation ener(cid:2) gies of the O2, N2, and NO molecules, should be taken into account.  The amount of the energy that is removed from the overheated regions also can be divided into two compo(cid:2) nents: one caused by the re(cid:2)radiation from relatively cold lateral surfaces qrad (qrerad, according to some authors) and due to the heat transfer deeper into the material by the mechanism of thermal conductivity qcond.  The re(cid:2)radiative component (qrerad) is a function of the ambient and wall temperatures, of the Stephan- Boltzmann coefficient σ, and of the emissivity ε, which depends on the temperature and can vary from 0 to 1 (ε = 1 for the blackbody).  The conductive heat transfer (qcond) depends on the local temperature gradient and on the thermal con(cid:2) ductivity of the material (λ).  Generalizing the above(cid:2)said, we should separate some characteristics of the materials, which would make it possible to select them as the most promising for the use in the hypersonic flight vehicles: the mate(cid:2) rials must have, besides high melting points and phase stability in a wide temperature range, also a compara(cid:2) tively high oxidation resistance (in particular, in reac(cid:2) tions with atomic oxygen); minimum catalytic activity in the exothermic reactions of surface recombination; high thermal conductivity, due to which heat can be removed from the strongly overheated regions; and high emissivity for improving the process of the re(cid:2) radiation of the heat obtained from the relatively cold lateral surfaces.  2. ON THE SPECIFIC FEATURES OF THE OXIDATION OF ULTRA(cid:2)HIGH(cid:2) TEMPERATURE MATERIALS BASED  ON BORIDES OF METALS  The processes of oxidation of materials based on refractory borides, first of all, zirconium and hafnium borides with the addition of silicon carbide, have been described in much detail and exhaustively in [2, 8, 10, 35, 36, 38-42, 49, 52-55, 75-77, 87, 89-94, 103, 107, 111, 118, 178, 179]; therefore, we will not con(cid:2) sider in detail this question, but only briefly touch some crucial points. As is known, the oxidation of materials such as Zr(Hf)B2-SiC at temperatures lower than ~1000- 1200°C (depending on the composition of the atmo(cid:2) sphere and pressure) occurs according to the following reaction:  Z rB 2  c(  )  +  5/2O 2  g(  )  =  Z rO 2 c(  )  +  B 2O 3 l(  ) .  (1)  At temperatures higher than 450°C, the boron oxide is liquid; due to the large contact(cid:2)wetting angle on the zirconium and hafnium oxides, it is distributed in the available pores, filling them and hampering the further diffusion of oxygen  into  the material. An increase in the temperature to ~1100°C significantly decreases the protective properties of the layer of liq(cid:2) uid B2O3 because of its active evaporation from the surface of the sample:  B 2O 3  l(  )  =  B 2O 3 g(  ) .  (2)  To solve this problem, silicon carbide is added into the composition of the high(cid:2)temperature material; at temperatures higher than 1000-1200°C (depending on SiC dispersity, pressure, and the composition of the atmosphere), it also undergoes oxidation:  SiC c(  )  +  3/2O 2 g(  )  =  SiO 2 l(  )  +  CO g(  ) .  (3)  The evolving silicon dioxide forms with the boron oxide a viscous borosilicate glass with a high boiling point, which hampers oxygen diffusion to the nonoxi(cid:2) dized  regions of  the material  (according  to  the Stokes-Einstein equation, the diffusion coefficient is inversely proportional to the viscosity of the medium through which the transport takes place). This pre(cid:2) vents the development of the process of further oxida(cid:2) tion and, also, leads to the suppression of the transi(cid:2) tion of B2O3 into the gas phase because of the decrease in its activity a in the SiO2-B2O3 system. Figure 8 [45] shows the results of the investigation (performed at the Kármán Institute for Fluid Dynam(cid:2) ics, Belgium) of the behavior of a ZrB2-SiC sample; as can be seen, with the formation at ~1400-1550°C of a viscous SiO2(cid:2)containing melt depleted of B2O3 due to its evaporation (whose formation in the course of the subsequent oxidation of ZrB2 is hampered by the barrier of borosilicate glass), the passage of the boron(cid:2)containing species into the gas phase practi(cid:2) cally ceases. The majority of researchers note that upon the oxi(cid:2) dation of Zr(Hf)B2-SiC materials there is formed a  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013    \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1675  Temperature, °C 1700  1500  1300  1100  .  u  .  a  ,  y  t  i  s  n  e  t  n  I  (cid:3)  g  i  S  w  a  R  .  u  .  a  ,  l  a  n  BO2  BO  B  Sample 7 Sample 17  λ = 518.8 nm  λ = 404.1 nm  λ = 249.9 nm  0  50  100  150  200  250  300  350  Time, s  Fig. 8. Variation of the intensity of lines in the emission spectrum, which corresponds to various forms of the gas(cid:2) eous compounds of boron in the near(cid:2)surface zone of the ZrB2-30 vol % SiC sample in the course of the investiga(cid:2) tion of its behavior in the high(cid:2)enthalpy flows of dissoci(cid:2) ated air using the plasmatron of the Kármán Institute for Fluid Dynamics, Belgium [45]; the upper curve reflects the temperature conditions.  SiO2  Depleted of SiC  (а)  10 μm  ZrB2-SiC  (b)  10 μm  SiO2  Depleted  of SiC  10 μm  ZrB2-SiC  (d)  10 μm  laminar structure of the oxidized region with the boro(cid:2) silicate glass predominantly concentrating on the sur(cid:2) face of the material; below, there is located a layer of porous ZrO2 or HfO2 partially filled with glass; further, in  certain  cases  there  is  observed  a  layer  of ZrB2(HfB2), having a porous structure and depleted of SiC; beneath, an unoxidized material is located (Fig. 9) [53]. The situation is explained in [8, 53, 55] and sub(cid:2) sequent works [35, 36, 42]. As a result of the formation of diffusion difficulties for the transport of oxygen into the material bulk on the interphase boundary, there is created a reduced partial oxygen pressure, which is insufficient  for  the oxidation of ZrB2(HfB2), but which ensures at elevated temperatures (>1500°C) the passage  to  the active oxidation of silicon carbide according to the reaction  S iC c(  )  +  O 2 g(  )  =  SiO g(  )  +  CO g(  ) .  (4)  The oxidized ZrB2(HfB2) creates a porous carcass(cid:2) like  structure  filled with  liquid borosilicate glass, which begins extruding outside due to the arising gas(cid:2) eous products, but is not yet torn off from the surface because of the adhesion by capillary forces with the walls of thin channels in ZrO2 or HfO2. Thus, in the porous (often, with cylindrical pores, although by no means always) massif of  refractory oxides, which ensures mechanical strength to the oxidized layer (the so(cid:2)called “solid pillars,” additionally strengthened due to the interaction with the liquid glass, which is partially located in the pores and partially covers the carcass and forms the so(cid:2)called “liquid roof”), there are formed cavities filled with gases, e.g., the products of the active oxidation of silicon carbide (SiO and CO) diffusing from the zone of reaction to the surface, and oxygen, which is transported through the borosilicate glass and reacts on the lower boundary of glass with the silicon monooxide, thereby increasing the thickness of the liquid layer:  S iO g(  )  +  1/2O 2  g(  )  =  SiO 2  l(  ) .  (5)  (c)  Schematically, this process is shown vividly in Fig. 10 [35]. Such a distribution of layers is sufficiently advan(cid:2) tageous from the viewpoint of heat transfer upon the aerodynamic heating, since it is known that the cata(cid:2) lytic activity of glassy coatings in the reactions of sur(cid:2) face recombination is substantially lower than that of the zirconium or hafnium oxides, and the emissivity is determined in a large measure by the materials located under the layer of glass (Zr(Hf)B2, SiC), which have ε ~ 0.8-0.9 at different temperatures. At the same time, for ZrO2 and HfO2 the emissivity in the high(cid:2) temperature region can be equal to 0.5-0.6. However, an increase in the temperature connected with a given increase in the power of the incident heat flux leads to the surface evaporation of not only boron(cid:2) containing compounds but also of silicon(cid:2)containing ones (SiO, SiO2), which can clearly be seen in Fig. 11 [8, 10]: when the temperature of the surface of the model sample of composition ZrB2-30 vol % SiC exceeds 1700-1750°C (second step in the heating  Fig. 9. (a) Microstructure of a polished section of a ZrB2- 30 vol % SiC sample and (b, c, d) the distribution maps of (b) O, (c) Si, and (d) Zr after the sample oxidation at 1500°C for 30 min [53].  curve), there occurs a sharp increase in the amount of boron(cid:2)containing molecules; in addition, there also appear signals from silicon(cid:2)containing compounds. Dur(cid:2) ing the entire highest(cid:2)temperature stage (for ~1 min), there is maintained a high concentration of silicon and boron oxides in the gas phase, which means their evap(cid:2) oration from the glassy surface layer, and the appear(cid:2) ance of a porous carcass of ZrO2 on the surface (ear(cid:2) lier, it was located below) with completely different catalytic and emissive characteristics, which also leads to a change in the chemistry of the surface. A similar situation was observed in [56], where after treatment of the sample of HfB2-35 vol % SiC in the flow of dis(cid:2) sociated air at a temperature of 2600-2650°C for  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013        \\x0c', '1676  SIMONENKO et al.  l12  l23  1  2  l3a  3  a  MeB2 + 5/2 O2 = MeO2 + B2O3  1(cid:3)fs  MeB2  MeO2  O  i  S  2  -  B  fs  SiC  O2(dissolved)  SiO/CO, CO2 B2O3-SiO2  Boundary layer diffusion SiO2(g) SiO(g) B2O3(g)  i1  i2  i3  a  CO (assumed not(cid:3)rate limiting)  SiC + 2CO2 = SiO + 3CO  SiO + 1/2O2 = SiO2 2CO + O2 = 2CO2  B2O3(l) = B2O3(g) SiO2(l) = SiO2(g) SiO2(l) = SiO(g)  Fig. 10. Formation of the porous carcass of ZrO2, the upper layer of the liquid glass adhered with the pores in ZrO2 by capillary forces, and the zone depleted of SiC because of its active oxidation (schematic) [35].  9 min there was formed a  laminar oxidized zone, which can be conditionally  subdivided  into  four regions (see Fig. 12); the upper layer, according to the EDX analysis, consists almost exclusively of HfO2. The emissive capacity of Zr2-SiC measured in [10] under hard heating conditions decreased only from ~0.9 to ~0.8, whereas the catalytic activity of the sur(cid:2) face changed radically: from ~3 × 10-3 to 0.2 -0.3. This phenomenon is indirectly confirmed by the data of [56], since the absence of the explicit dependence of the surface temperature on the external pressure of dissociated air tells precisely about the formation of a highly catalytic upper layer. In [10], on the surface there were also observed particles of zirconium diox(cid:2) ide, which, probably, leads to an increase in the chem(cid:2) ical component of heat fluxes and, correspondingly, to an increase in the surface temperature, which can lead to the destruction of the sharp edges of articles.  The partial oxygen pressure in the gas mixture of a high(cid:2)enthalpy flow exerts an essential effect also on the character of the oxidation of composite materials Zr(Hf)B2-SiC [42, 52]. In [52], there was performed a systematic study of the oxidation resistance of the hot(cid:2)pressed samples of the ceramic ZrB2-20 vol % SiC in the microwave(cid:2)discharge plasma apparatus at the temperature of 1500°C and different partial pres(cid:2) sures of oxygen (60 000, 20 000, 4000, 1000, 100, and 10 Pa). It is shown that at high pressures of O2 (60 kPa) there is formed an almost smooth glassy layer (accord(cid:2) ing to the XRD data, the crystalline phase, which is contained  in  its volume,  is monoclinic ZrO2); at p(O2) = 20 kPa and, especially, at 4 kPa, above the sur(cid:2) face there begin appearing particles of ZrO2 (Fig. 13) due to the evaporation of SiO2 and B2O3. At the lower partial pressures of oxygen (1000, 100, 10 Pa), there is formed only a porous ZrO2 carcass; SiC in this case undergoes an active oxidation (this fact is confirmed by the SEM micrographs of the sections of the samples [52]). A change of the thickness of the oxidized region depending on p(O2) has an extreme nature. Thus, for predicting the behavior of oxidation processes on the surface and in the bulk of Zr(Hf)B2-SiC materials, it  is extremely important to well understand the thermo(cid:2) dynamic and kinetic aspects of the passage from the passive to active oxidation of SiC [21-25], since in some experimental works [26, 42] it is shown that in the oxidized layer there are also present carbon inclu(cid:2) sions, which, according to the thermodynamic calcu(cid:2) lations, can form at even lower partial pressures of O2 according to the following equations:  S iC c(  )  +  SiO 2  (  s l,  )  =  2 SiO g(  )  +  C s(  ) ,  S iC c(  )  +  1/2O 2  g(  )  =  SiO g(  )  +  C s(  ) .  (6)  (7)  On the whole, with increasing distance to the sur(cid:2) face of the material there appears an essential differ(cid:2) ence in the compositions of the solid phase and gas(cid:2) eous reagents, as well in the partial pressures of differ(cid:2) ent components and temperatures (as a result of the low heat conductivity of ZrO2 and HfO2); therefore, it is necessary to examine all possible reactions in this complex system, even those that are very improbable, at first glance, including processes with the participa(cid:2) tion of elementary silicon, oxycarbides and oxyborides of silicon and zirconium (hafnium), or the formation of zircon and hafnon. Since the chemical properties of the components of ultra(cid:2)high(cid:2)temperature ceramics do not permit one to completely avoid the oxidation processes and degra(cid:2) dation, active and diverse works on searching methods that could maximally level(cid:2)off these negative changes are carried out. The overwhelming majority of works in this field is directed on the creation of dense ceramic materials, with a zero porosity if possible, which not only opti(cid:2) mizes the mechanical characteristics of the starting material Zr(Hf)B2-SiC (or materials on the base of super(cid:2)refractory carbides), but also hampers the pro(cid:2) cess of the penetration of oxygen into the bulk of the UHTC. The following methods are used most fre(cid:2) quently for these purposes: (1) cold pressing with subsequent sintering [1, 10, 84, 121, 129] (in this case, obviously, it is impossible avoid the application of special sintering components, e.g., Si3N4, MoSi2);  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1677  MeB2  MeO2  SiC  SiO/CO, CO2 B2O3-SiO2  O2(diss.)  i1  i2  i3i  a  Evaporation limited by diffusion through porous channels  SiO2(g) SiO(g) B2O3(g)  Sample 1.5  2700  2400  2100  1800  1500  10 000  B(cid:3)atom (249.92 nm)  K  ,  e  r  u  t  a  r  e  p  m  e  T  1000  .  u  .  a  10 000  Si(cid:3)atom (288.16 nm)  ,  y  t  i  s  n  e  t  n  i  l  a  n  g  i  s  1000  10 000  W(cid:3)atom (400.99 nm)  w  a  R  1000  0  1  2  3  4 5 6 7 Test time, min  8  9  10  Fig. 11. (a) Schematic representation of the distribution of layers and the processes that occur at the surface as a result of the aerodynamic heating to above 1750°C [35]; and (b) the results of the action of high(cid:2)enthalpy flows of dissoci(cid:2) ated air on the sample of ZrB2-30 vol % SiC [10]: the upper curve reflects the temperature regime of treatment; the lower curves show changes in the content of Si(cid:2) and B(cid:2) containing molecules in the gas phase over the surface of the sample (according to EDS data).  (2) hot pressing, predominantly at temperatures of 1900-2100°C and loads of ~20-50 MPa (it is used in the overwhelming majority of works [5-9, 14-17, 19, 26, 30, 32, 34, 37, 38, 40, 42, 45, 47, 49, 52, 65-73, 75, 76, 83-87, 89, 90, 92, 93, 95, 106, 107, 117, 118, 170, 187]);  (3) spark plasma sintering (SPS) [31, 43, 54, 56, 95-102, 114, 174, 175, 184]; this is the most progres(cid:2) sive method, which makes it possible to obtain dense samples of even most refractory carbides at lower tem(cid:2) peratures [112, 113, 183, 185, 186] (which can be obtained in a bulk nanostructured state if nanosized  300 μm  Fig. 12. Morphology of the polished section of an HfB2- 35 vol % SiC sample after plasma(cid:2)chemical treatment, according to the data of optical microscopy (central zone of the sample) [56].  initial powders are used for their preparation, whose synthesis is by itself a complex chemical(cid:2)engineering problem [188-191]).  One more method of obtaining ultra(cid:2)high(cid:2)temper(cid:2) ature ceramics is based on different modifications of the  high(cid:2)temperature  self(cid:2)propagating  synthesis (SHS), which is initiated by the application of an ele(cid:2) vated temperature and external pressure, or reaction sintering [44, 61, 81, 82, 103, 104]. As the starting sub(cid:2) stances,  there are used powders of zirconium or hafnium,  their hydrides,  silicon,  silicon carbide, oxides of zirconium or hafnium in reactions with B4C, and other variations.  Furthermore, methods of polymer impregnation and pyrolysis (PIP) are employed, although on a sub(cid:2) stantially smaller scale (mainly for creating refractory protective matrices and coatings on Cf/C and Cf/SiC composites), as well as chemical vapor deposition (CVD)  and  often  also  combinations  of  these approaches [3, 4, 58, 62, 63, 108-110, 119, 172, 173, 180]. For the efficient application of these methods, the development of precursors and starting reagents with given properties is necessary [181, 182].  An approach to an increase in the oxidation resis(cid:2) tance via the formation of maximally dense materials is  reasonable  (especially,  taking  into account  the requirements  to  the mechanical  characteristics). However, in [56], where we studied the behavior of HfB2-35 vol % SiC samples with a high porosity (~28-29%) under the action of a flow of dissociated air with the application of an induction plasmatron at temperatures of more than 2600°C (for 9 min), the samples successfully withstood under such severe con(cid:2) ditions; the thickness of the oxidized layer was 600- 650 μm, and a 30(cid:2)min treatment led to a 2% increase in mass rather than to an ablation due to, e.g., the  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013            \\x0c', '1678  SIMONENKO et al.  (а)  (c)  (e)  100 μm  (b)  100 μm  100 μm  (d)  5 μm  5 μm  (f)  5 μm  Fig. 13. Microstructure of the surface of ZrB2-20 vol % SiC samples after oxidation at 1500°C (for 30 min) at dif(cid:2) ferent partial oxygen pressures: (a) 60 000, (b) 20 000, (c) 4000, (d) 1000, (e) 100, and (f) 10 Pa [52].  destruction of the porous upper layer of HfO2 and intensive evaporation of borosilicate glass. It is possi(cid:2) ble that the absence of essential mechanical damages of the sample is connected with its geometry in the form of a tablet with a diameter of 15 mm, which is nearer to the model of “blunt” rather than “sharp” edge being subjected to a maximum load; however, it cannot be excluded also that the low porosity by no means necessarily leads to the best results.  A change in the content of SiC, naturally, should strongly affect the oxidation resistance of UHTCs, since the product of its oxidation is SiO2, which forms  Table 2. Thickness of layers that are formed as a result of oxidation of ZrB2-xSiC materials (x = 20, 35, 50, 65, and 80 vol %) at 1773 K for 50 h [106]  Designation of the sample  Thickness  of SiO2(cid:2)based  layer, μm  Thickness  of the SiC(cid:2) depleted layer,  μm  S80Z20 (ZrB2-80 vol % SiC)  S65Z35 (ZrB2-65 vol % SiC)  S50Z50 (ZrB2-50 vol % SiC)  S35Z65 (ZrB2-35 vol  % SiC)  S20Z80 (ZrB2-80 vol  % SiC)  490 ± 50  65 ± 10  100 ± 10  120 ± 10  120 ± 10  absent  absent  80 ± 20  170 ± 10  260 ± 75  a diffusion barrier for oxygen—a borosilicate glass. It has been noted that the use of nanosized powder of sil(cid:2) icon carbide [29, 102] makes it possible both to slow down the occurrence of oxidation processes in com(cid:2) parison with its coarse(cid:2)crystalline state and to improve the mechanical characteristics of materials. In the majority of literature sources, the recommended con(cid:2) tent is from 10 to 30 vol % SiC. In a number of works, it is recommended to maximally decrease the fraction of silicon carbide, since the effect of the formation of SiC(cid:2)depleted zones can arise in porous ZrB2 or HfB2 because of an active oxidation of SiC in the bulk of the material. However, both the data of [56] on the tests of the of HfB2-35 vol % SiC sample and the system(cid:2) atic investigation [106] of the oxidation behavior of ZrB2(cid:2)based materials, which contain 20, 35, 50, 65, and 80 vol % SiC, give grounds to doubt the complete correctness of this thesis. Thus, for ZrB2-xSiC sam(cid:2) ples obtained under identical conditions by the method of hot pressing, the study of the oxidation process upon isothermal holding at a temperature of 1773 K for 50 h has shown (Table 2) that, in proportion to an increase in the content of SiC (except for the ZrB2-80 vol % SiC sample), there is observed a tendency toward a reduction in the thickness of both the layer of the borosilicate glass and the SiC(cid:2)depleted region (for the samples with x = 65 and 80 vol % SiC, its formation was not  fixed at all). A similar picture also was observed in the course of experiments on the oxidation of materials at 2073 K for 20 min. Therefore, the authors of [106] have concluded that the content of 10-30 vol % silicon carbide in the composition of a UHTC based on ZrB2 or HfB2 is by no means neces(cid:2) sarily optimum from the viewpoint of oxidation resis(cid:2) tance.  Another approach to an increase in the resistance of materials  to  the action of both molecular and atomic oxygen is the application of alternative alloying additives, whose oxidation will lead to the formation of silicon dioxide and, as a result, of a borosilicate glass. As an example of such compounds, there can be mentioned silicon nitride [40, 84], many silicides, e.g., molybdenum [1, 48, 111, 115, 117], zirconium [91], and tantalum or niobium silicides [15, 32, 89, 95, 100]. The latter metals form (except for SiO2) also compar(cid:2) atively refractory oxides, which, being dissolved in the borosilicate glass, must increase its viscosity at ele(cid:2) vated temperatures, thereby decreasing losses due to the ablation of the liquid phase and evaporation of its components.  An original concept for increasing oxidation resis(cid:2) tance by the directional modification of the glasslike layer was advanced by the authors of [2]. They sug(cid:2) gested introducing additives such as VB2, NbB2, TaB2, TiB2 or CrB2 into the initial powder, although the oxi(cid:2) dizing stability of these compounds is significantly inferior to that of ZrB2 or HfB2. The essence of this method is in that the introduction of oxides of the IV- VI Group transition metals into the borate and silicate  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1679  glasses leads to phase separation, which, in turn, leads to a noticeable increase in the temperatures of the liq(cid:2) uidus curve and an increase in the viscosity of the sys(cid:2) tem, and, as a result, to a reduction in the rate of oxy(cid:2) gen diffusion and suppression of processes of evapora(cid:2) tion from the B2O3 glass. The degree of the phase separation of glassy com(cid:2) positions in the upper protective layer is higher, the greater the z/r2 ratio, where z is the degree of oxida(cid:2) tion, and r is the ionic radius of the transition metal. For all the metals that enter into the composition of the above(cid:2)mentioned diborides, this ratio is higher than that for the zirconium or hafnium cations; there(cid:2) fore, their introduction will favor the phase separation of the borosilicate glass that is formed on the surface of the ceramics, for example, of ZrB2-SiC [2]. The modification was performed via partial substitution of the corresponding diboride for ZrB2 (2-20 mol %) with the fixation of the content of silicon carbide. The relative efficiency was determined using the isother(cid:2) mal holding of the samples (obtained under identical conditions) at temperatures of 1200-1600°C in an oxygen-argon gas mixture imitating air. For the esti(cid:2) mation, the changes in the mass due to the oxidation were determined; in addition, the thickness of the oxi(cid:2) dized layer was estimated, and the microstructure and the composition of the surface and of the transverse cut of the samples were studied. The experiment at a temperature of 1300°C showed that, among the sam(cid:2) ples that contain 10 mol % CrB2, NbB2, TaB2, TiB2, and VB2,  the  smallest  increase  in  the mass was observed for the ceramic material containing TaB2; the thickness of the oxidized layer was less than half of that corresponding to the base material ZrB2-SiC. The sequence of the mass increments for the samples with different alloying components corresponded to the sequence of decrease in the z/r2 ratio for the Ta+5, V+4, Nb+5, Ti+4, and Cr+3 ions. The SEM study of the sur(cid:2) face confirmed the presence of phase separation; in the partially crystallized glasslike matrix (enriched in zirconium and tantalum), large inclusions of borosili(cid:2) cate glass are observed. The XRD analysis showed that there are crystalline phases of two types: ZrO2 and a small amount of TaZr2.75O8. The opportunity of con(cid:2) trolling the oxidation resistance of ZrB2 via the cre(cid:2) ation of conditions for the phase separation in the pro(cid:2) tective surface glassy layer was successfully shown on the example of CrB2, TiB2, TaB2, NbB2, ZrB2-Si3N4, Ti3SiC2, and Si3N4 ceramics [2]. On the whole, upon the introduction of borides of IV-VIB(cid:2)Group metals, a certain improvement in the oxidation(cid:2)resistance of UHTCs [10, 73, 80, 118] is observed; the same reasons appear to cause the presence of the separation of glass upon the introduction of TaSi2 and NbSi2 into the initial charge and, in certain cases, the modification by the tantalum carbide; however, it should be remembered that the oxidation behavior of samples in the furnaces differs significantly from that under the conditions of aerodynamic heating.  ) s  t  n  u o  C  (  m(cid:3)ZrO2  t(cid:3)ZrO2  La2Zr2O7  y  t  i  s  n  e  t  n  I  20  30  40  50  60  70  2θ, deg  Fig. 14. X(cid:2)ray diffraction pattern of the surface layer of a ZrB2-20 vol % SiC-10 vol % LaB6 sample.  One of the essential problems of Zr(Hf)B2-SiC materials is the fact that, as a result of their oxidation, there are formed oxides that are very refractory and have a vapor pressure that is low up to super(cid:2)high tem(cid:2) peratures but which, simultaneously, are characterized by phase transitions that are accompanied by a large volume change and by linear thermal expansion coef(cid:2) ficients that differ strongly from those characteristic of the compact part of the samples. Therefore, in the majority of cases upon the study of the microstructure of metallographic sections after oxidizing tests there is observed spalling of the oxidized part of the sample from the base. The process of spalling occurs as early as at the stage of the tests that simulates the high(cid:2) enthalpy air flow. In order to stabilize ZrO2 in either the tetragonal or cubic modifications or in the compo(cid:2) sition of a compound that exhibits phase stability in the entire range of temperatures from room tempera(cid:2) ture to the tentative operating conditions, compounds of rare(cid:2)earth elements (most frequently, of lanthanum and gadolinium in the form of Ln2O3 [51, 83] or LaB6 [19, 51]) are introduced into the starting powders. Testing  in  the  flame  of  an  oxyacetylene  torch (~2400°C) showed an improvement of the adhesion of the oxidized layer in the bulk part of the sample; the XRD indicates (Fig. 14 [19]) that, along with the for(cid:2) mation of the monoclinic modification of ZrO2, there is also formed a tetragonal phase and cubic La2Zr2O7.  In [32], the authors introduced 5 vol % metallic iri(cid:2) dium (particles with dimensions less than 44 μm) into the charge (HfB2-10 vol % SiC-5 vol % TaSi2). For the samples obtained at 1800-1900°C,  there was observed a decrease in the thickness of the oxidized layers; it was noted that this metal is not dissolved in HfB2 and SiC; therefore, it is concentrated at grain boundaries and hampers the diffusion of oxygen.  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013    \\x0c', \"1680  SIMONENKO et al.  5 3 5  .  6 6  Internal design schematic  STA 0.000  STA 5.047  HfB2/SiC Nosetip  R  4 1 7  .  0 1  + X  STA 4.496 STA 4.589 Tier 1 Sensor TC4 and TC5  ZrB2/SiC/C Ph  STA 7.560 Tier 2 Sensor TC1, TC2, and TC3  MK12A Nose  STA 12.040  Fig. 15. Location of a sample of the UHTC material HfB2-SiC in the head of an Mk12A ballistic rocket Mi(cid:2) nuteman III (launch SHARP(cid:2)B2) (schematic) [27].  Fig. 16. Interaction between the sample placed into a holder and the flow of dissociated air (electric(cid:2)arc plasma(cid:2) tron) [45].  3. TESTS OF UHTC MATERIALS  THAT SIMULATE THE RE(cid:2)ENTRY  CONDITIONS FOR THE SUPERSONIC  FLIGHT VEHICLES  As was already said, the selection of a material for the use in the special systems of thermal protection is determined  by  the  construction, medium,  and requirements imposed on the properties of the mate(cid:2) rial. The final choice of the material must be based on the tests of the finished material in the appropriate medium; the tests carried out in an incorrect medium or under incorrectly selected conditions can lead to the rejection of a promising material or to the selection of an inappropriate material.  3.1. Flight Tests  At the end of the 1990's, the UHT composites have been used in two flight tests, SHARP(cid:2)B1 and SHARP(cid:2)B2, which were conducted at the Ames NASA Research Center [20, 27]. The first flight test (SHARP(cid:2)B1) took place on May 21, 1997; the second (SHARP(cid:2)B2), on September 28, 2000. The purpose of the experiment was to confirm the possibility of using blunted nose tips of a small radius (3.6 mm) and of sharp leading edges without using systems of active cooling. In the SHARP(cid:2)B2 experiment, the element made of the HfB2-SiC material was mounted in the nose part of the conical nose block Mk12A of the intercontinental ballistic missile Minuteman III (Fig. 15). The temper(cid:2) ature of the sample surface was recorded with the aid of thermocouples. The beginning of the ablation of the nose tip was registered at the height of 58 km at the surface temperature equal to 2760°C. After the com(cid:2) pletion of the flight tests, efforts have been undertaken to modify the UHTC and eliminate the deficiencies observed during the SHARP B2 tests; the materials were found to be characterized by a very low Weibull modulus (~4) and a low strength because of the pres(cid:2) ence of treatment(cid:2)induced defects [20, 27].  3.2. Ground Tests  On the whole, the flight tests are extremely expen(cid:2) sive and organizationally complex;  therefore,  for searching promising materials  there are predomi(cid:2) nantly used ground(cid:2)based installations, which simu(cid:2) late aerodynamic loads that can exist in a flight on pre(cid:2) determined trajectories with required characteristics. The simplest variant that permits comparison between the proposed materials seems to be the use of the flame of an oxyacetylene torch. More correct and more flexi(cid:2) ble variant is the use of arc(cid:2)jet plasmatrons [1, 6, 7, 10, 11, 15-18, 26, 30, 32, 38, 40, 45, 46, 48, 49] (Fig. 16 [45]), which makes it possible to change the rate of flow, the pressure in the chamber, and some other parameters in relatively wide limits.  Owing to the conducted plasma(cid:2)chemical investi(cid:2) gations of Zr(Hf)B2-SiC materials, in a number of cases it was possible to determine the emissivity coef(cid:2) ficient (ε) and catalytic(cid:2)recombination coefficient (γ) under various treatment conditions, and also to reveal their changes depending on the surface modification in the course of oxidation and increase in temperature.  Below, we consider the most interesting and signif(cid:2) icant, in our opinion, works concerning the testing of the some UHT materials based on ZrB2 and HfB2 under the action of flows of dissociated air.  3.2.1. Tests, in which the samples have almost flat front surface [6, 7, 10, 15, 32, 45, 46, 48, 49, 56].  Gasch et al. [6] have studied the behavior of hot(cid:2) pressed samples of HfB2-20 vol % SiC under the effect of a flow of dissociated air using an electric(cid:2)arc plasmatron. The samples were made as details with a  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013    \\x0c\", 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1681  flat surface, which were mounted in graphite holders coated by SiC (Fig. 17). The studies were carried out using two regimes: with a comparatively mild action (thermal flow Q ~ 285 W/cm2; pressure 0.005 MPa) and under somewhat harder conditions (Q = 350 W/cm2, pressure 0.007 MPa). In each test, the samples were placed into the plasma jet doubly and held for 10 min each time.  It has been found [6] that, when using the “soft” regime, upon both the first and the second introduc(cid:2) tion of the sample into the flow of dissociated air, a temperature of 1690°C was established on the surface of the sample and remained almost constant during the holding. For the sample subjected to the severer treatment, during the first treatment there is estab(cid:2) lished a temperature of ~1800°C, which in the course of the test grows at an increased rate and at the end of the holding time becomes more than 2200°C. During the repeated introduction of the sample into the jet, a temperature of ~2400°C  is established sufficiently rapidly (in ~200 s). The SEM data made it possible to determine the thickness of the oxide layer and of the SiC(cid:2)depleted zone: for the sample subjected to the “soft” action, they are 70 and 2 μm, respectively; for the sample tested under the severer conditions, 340 and 740 μm, respectively.  The authors of [6] make the conclusion that in the first case at temperatures approaching 1700°C and at a low pressure the oxidation of the components makes it possible to create a passivating layer, while in the sec(cid:2) ond case there occurs a passage to the active oxidation of SiC. With  time,  the protective glassy  layer  is removed away from the surface of the sample with the formation of pores in the form of channels, stimulat(cid:2) ing active oxidation of SiC with the formation of gas(cid:2) phase products; as a result, there occurs an increase in the thickness of the layer depleted of this component. An increase in the temperature in the second half of the first treatment for this sample is explained by the authors of [6] by the effect of a combination of several factors, in particular, by a gradual removal of the ini(cid:2) tially formed SiO2-B2O3(cid:2)based layer and the emer(cid:2) gence of HfO2 onto the surface, which leads to an increase in the catalytic activity and a reduction of the emissivity; an increase in the temperature can also occur because of the  low thermal conductivity of HfO2, which prevents the removed of heat by the con(cid:2) ductive mechanism.  In [7] the hot(cid:2)pressed dense samples of ZrB2- 20 vol % SiC (in the form of bodies with a flat front surface) were tested using an electric(cid:2)arc plasmatron at  subsonic  regimes and heat  fluxes of 1.7 and 5.4 MW/m2 (average entalpiy H0 = 2-4 MJ/kg; time of treatment, 10 min). At the heat flux of 1.7 MW/m2, the temperature of the surface of a flat sample was 1650°C and no mass change was fixed. In the second case (Q = 5.4 MW/m2), there occurred an active oxi(cid:2) dation and destruction of the material (in fact, a froth(cid:2) ing of the material, see Fig. 18); the loss of mass was  Water cooled sting  Coated graphite model holder  Arc jet flow  Ames UHTC model  Fig. 17. The external appearance of a “sample-holder” system during plasma(cid:2)chemical tests [6].  (а)  (b)  (c)  Fig. 18. The external appearance of a sample of composi(cid:2) tion ZrB2-20 vol % SiC (a) during treatment and (b, c) after plasma(cid:2)chemical  treatment using a heat  flux of 5.4 MW/m2 (surface temperature: 2300°C) [7].  15.75%, and a change in the thickness, ~3 mm; the maximum  temperature  of  the  surface Tmax has exceeded 2300°C. As a result of the ablation of a large amount of molten oxidation products by the flow and of the action of high shear loads, the edges of the tablet were  subjected  to  a  larger  erosion;  they were smoothed. According to the SEM data, a dense layer almost completely consisting of ZrO2 with  small inclusions of SiO2 was formed at the surface; further, a more porous layer of ZrO2 was located, which then passed into a zone depleted of SiC (25 μm).  From the above(cid:2)described material [7], samples with a sharp leading edge (with the radius of curvature ~3.5 mm) were prepared, which were subjected to the action of dissociated air (Mach number, 2.7; time of action, 10 min). As a result, the temperature grew sharply (to 1350°C) and then established at a level of 1450°C. No crack formation occurred, and the mass change was 0.03%, whereas  the C/SiC(cid:2)composite detail of a similar morphology tested under the same conditions has lost 8.3% of its mass, and a substantial part of its leading edge became destroyed (Fig. 19). It has been shown that the ZrB2(cid:2)based material is more promising as the material for the supersonic applica(cid:2) tions as compared to the classical C/SiC composite.  The authors of [10] investigated the phenomenon of a jumplike growth of temperature on the surface of materials such as ZrB2-30 vol % SiC, (Zr0.96W0.04)B2,  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1682  SIMONENKO et al.  (а)  (c)  (b)  (d)  Fig. 19. Photographs of samples with a sharp leading edge: (a, c) before and (b, d) after plasma(cid:2)chemical treatment for (a, b) ZrB2-SiC and (c, d) C/SiC [7].  HfB2-20SiC (HP)  HfB2-20SiC (FAS)  HfB2-10SiC-5TaSi2 (HP)  HfB2-10SiC-5TaSi2 (FAS)  Did not arc(cid:3)jet best.   Sample cracked  durirg fabrication  HfB2-15SiC-5TaSi2-5Ir (HP)  Fig. 20. Photographs of the surface of model samples after electric(cid:2)arc testing [32].  and (Zr0.96W0.04)B2-30 vol % SiC, which were used to prepare samples of complex shape with the working  part in the form of a cylinder with a diameter of 30 mm  and a radius of curvature of the lateral surfaces equal to  7 mm; as a result, details of mushroom shape are formed. The regime of the plasmatron operation in this case was as follows: mass rate of air, 16 g/s; static pressure in the chamber, 10 kPa; enthalpy of the inci(cid:2) dent flow, from 10 to 28 kJ/g, which was controlled by the variation of the plasmatron power from 170 to 390 kW; as the power increased to more than 330 kW, there was observed a jumplike increase in the temperature by several hundreds of kelvins (with a substantial loss of mass) in the absence of changes in the operational characteristics of the apparatus, which, as the authors of [10] believe, is related to a change in the energy bal(cid:2) ance on the surface rather than to an increase in the enthalpy of the incoming flow. During the experiment, the coefficients of emission and catalytic activity were determined. It has been shown that for all the samples the emissivity before the switching into the higher(cid:2) energy regime was 0.9-0.95, and after the switching, 0.75-0.9 (~0.8, mostly). The recombination coeffi(cid:2) cients upon the treatment using the smaller power var(cid:2) ied between 3.1 × 10-3 and 6.3 × 10-3; after an increase in the power, they grew sharply (to 0.16-0.37). The authors of [10] explain this by a combined effect of an increase in qchem, a decrease in qrerad, and a change in the chemistry of the surface.  In [32], the authors have studied samples of an HfB2(cid:2)based UHTC obtained by hot pressing (HP) and an electric field(cid:2)assisted sintering (FAS) with an addi(cid:2) tion of 10-20 vol % SiC, 5 vol % TaSi2, and 5 vol % Ir. The samples prepared by the FAS technique are char(cid:2) acterized by a smaller (by a factor of 1.5-2) grain size in comparison with the HP samples and by a some(cid:2) what increased hardness.  The tests for the oxidation resistance of model samples (diameter 25.4 mm, total height 8 mm) with a flat surface were carried out under the conditions that simulate the medium existing around a supersonic flight vehicle upon  its re(cid:2)entry to the atmosphere (Qcold(cid:2)wall ~ 250 W/cm2, duration 5 min) with the aid of an arc(cid:2)jet plasma setup. Figure 20 displays the photo(cid:2) graphs of the sample surfaces after arc(cid:2)jet tests in heat fluxes of ~250-280 W/cm2 for 5 min. The parameters of testing are given in Table 3. For the HfB2-20 vol % SiC  samples,  the  temperature of  the  surface  is ~1690°C (HP) and ~1530°C (FAS). An analogous sit(cid:2) uation is observed for the samples with an addition of TaSi2. The surfaces of the samples with SiC and TaSi2 remain smooth in the process of testing and after cool(cid:2) ing. However, in the case of the iridium(cid:2)containing sample, embedments are formed on the surface in the course of testing. On the surface of the samples after testing there is present a layer of SiO2, under which there is located a region depleted of SiC, with porous grains of HfB2. On the whole, the oxidation resistance of the samples with additions of TaSi2 is higher as com(cid:2) pared to the HfB2-SiC ceramic. However, the results for the HP samples show that, although the addition of TaSi2 leads to a decrease in the thickness of the oxide layer,  the  thickness of  the SiC(cid:2)depleted  region  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1683  Table 3. Conditions of the arc(cid:2)jet tests [32] (Q is the heat flux; τ, the duration of treatment; ε, emissivity; and L, thickness)  Sample  Q, W/cm2  τ, s  T, °C  ε  after tests  Thickness of the oxi(cid:2) dized layer L, μm   Thickness of the SiC(cid:2) depleted region L, μm  HfB2-20SiC (HP)  HfB2-20SiC (FAS)  HfB2-10SiC-5TaSi2 (HP)  HfB2-10SiC-5TaSi2 (FAS)  HfB2-15SiC-5TaSi2-5Ir (HP)  280  250  250  250  250  600  600  600  600  600  1690  1530  1560  1515  1590  0.67  0.87  0.89  0.89  0.87   HP = hot sintering; FAS = electric(cid:2)field(cid:2)assisted sintering.  13  3  7  3  4  24  8  34  6  9  increases. It was noted that upon addition of TaSi2 and Ir the thickness of the oxide layer decreases approxi(cid:2) mately by a factor of 3. The samples prepared by the FAS technique have a lowered thickness of the oxide layer and a lowered thickness of the region depleted of SiC in comparison with the samples obtained by hot pressing. The authors of [32] believe that the introduc(cid:2) tion of SiC particles with small grain size into HfB2 ensures  an  increase  in  the  extension  of  the diboride/SiC interface and a decrease in the distance between SiC particles, which makes it possible to form a protective oxide layer on the surface in a shorter time (see also [33]).  In [56], samples of HfB2-35 vol % SiC and HfB2- 25 vol % SiC in the form of tablets with a diameter of ~15 mm and a thickness of ~5 mm obtained by the method of spark(cid:2)plasma sintering (porosity ~28%) were subjected to the action of a flow of dissociated air with  the application of a VGU(cid:2)4 high(cid:2)frequency induction plasmatron [56]. The temperature of the sample surface measured using a color pyrometer was more than 2000°C. In the course of plasma(cid:2)chemical experiment, no one of the four samples studied has been destructed; the change in their mass due to the oxidation and the removal of the evaporated products by the flow of the air plasma was ≤2.4%. The joint use of the data of optical microscopy, SEM, and X(cid:2)ray computer microtomography makes it possible to assert that the obtained samples with a content of SiC equal to 25 and 35 vol % have a sufficiently high thermal(cid:2) shock resistance, since at a sharp reduction in the tem(cid:2) perature of the sample surface from 2000-2600°C to ~1000°C during 3-5 s neither cracking nor phase sep(cid:2) aration occurred. In the case of the sample that was subjected to the most severe action, the surface tem(cid:2) perature was >2000°C upon treatment for 11 min and >2600°C upon treatment for 9 min. The SEM and EDX analysis of this sample suggest that the combina(cid:2) tion of the high temperature of the surface (~2600°C) and the action of a flow of dissociated air leads to a rapid evaporation of the components of the oxidation(cid:2) produced borosilicate glass from the surface layer. As a result,  the outer  layer consists predominantly of porous HfO2; therefore, the authors of [56] assume that the further oxidation is limited by the diffusion of  atomic and molecular oxygen (into HfB2) and of gas(cid:2) eous oxidation products (from HfB2)  through  the available channels in HfO2 and in the more deeply lying layer of the SiO2-B2O3 glass. It was established that for this sample the total thickness of the oxidized layer, including SiC(cid:2)depleted region, is 600-650 μm (Fig. 12).  3.2.2. Testing model samples with a radius of curva(cid:2) ture equal to 0.5-10 mm [1, 18, 26, 30, 38]. The  authors of [1] have employed the arc(cid:2)plasma method to study materials based on HfB2 and HfC sintered without pressure in a temperature range of 1950- 2400°C. The following materials were selected for these studies: HfB2-5 vol % MoSi2 (HB5) and HfC- 5 vol % MoSi2 (HC5). The sintered specimens had a hemispherical shape; the surfaces were ground to a roughness of 0.25 μm. The coefficients of thermal emissivity were 0.9  (1600-2000°C)  for HB5 and 0.7 (1800-2400°C) for HC5. The samples were sub(cid:2) jected to the action of steady(cid:2)state(cid:2)enthalpy flows using an arc(cid:2)jet plasma installation; the distance from the samples to the edge of the torch was 6 cm; the composition of the gas flow was 75% Ar-25% N2; the mass flow rate, 1.45 g/s; H0 was varied from 20 to 28 MJ/kg at the atmospheric pressure. At the outlet of the torch, the plasma, which contains argon, molecu(cid:2) lar nitrogen, and atomic nitrogen, expands through a nozzle (with a diameter of 5 mm) and comes into con(cid:2) tact with the ambient air, which leads to the dissocia(cid:2) tion of the atmospheric oxygen and formation of a mixture consisting of Ar, O2, N2, NO, O, and N.  Figures 21a and 21b show the dependence of the surface temperature on the testing time for the HB5 and HC5 samples, respectively. The maximum sur(cid:2) face  temperature  for  the HB5  sample  reached 1950°C (Fig. 21a). The calculated heat flux at the stagnation point was 5-8 MW/m2. At the highest tem(cid:2) perature, the emissivity was equal to approximately 0.9. When  testing  the  first hafnium(cid:2)carbide(cid:2)based sample (HC5(cid:2)I),  the surface  temperature reached 2050°C (Fig. 21b). In the case of the second sample (HC5(cid:2)II), the surface temperature reached 2400°C and remained at this level for ~4 min. The correspond(cid:2) ing calculated heat flux was ~10 MW/m2. In both tests,  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1684  SIMONENKO et al.  2000 1750  1500  1250  1000  750  500  250  C  °  ,  T  (a)  c  d  HB5  b  a  a b c d  H = 20 MJ/kg  H = 22 MJ/kg  H = 24 MJ/kg  H = 26 MJ/kg  0  60  120  180  Time, s  C  °  ,  T  2500 2250 2000 1750 1500 1250 1000 750 500 250  0  (b)  HC5(cid:3)II HC5(cid:3)I  c  b  a  d  a b c d  H = 20 MJ/kg  H = 22 MJ/kg  H = 24 MJ/kg  H = 26 MJ/kg  60  180 300 420 120 240 360 Time, s  Fig. 21. Time dependences of the temperature profiles in the course of electric(cid:2)arc tests for (a) HB5 and (b) HC5(cid:2)I and HC5(cid:2)II model samples, and the corresponding specific total enthalpies [1].  the measured emissivity coefficient at the maximum temperatures was ~0.7.  The analysis of the transverse section of the HB5 sample [1] revealed that the oxidized layer consists of several sublayers and adheres well to the bulk of the sample, which means  the absence of micro(cid:2) and macro(cid:2)separation. Under the surface layer of SiO2 (~10 μm) there are located mainly large grains of HfO2 and particles, whose composition is close to SiO2. Inside the glassy phase, an insignificant amount of aluminum inclusions is located. The innermost layer of the oxidized part of the sample contains the oxides of molybdenum,  silicon, and hafnium; no  layer depleted of silicon was observed.  For the HC5(cid:2)I sample, the color of the surface changed after tests from dark gray to bright; on the sur(cid:2) face, only cracked HfO2 without glassy phase  is present. The SEM study of the section of this sample has shown that there is formed a laminar oxidized layer ~90 μm thick closely adhering to the bulk part of the material. The outer layer (directly adjoining the sur(cid:2) face) consists of HfO2 and of a SiO2(cid:2)based glassy phase, which partially fills the available pores. The intermediate layer contains a small amount of pores and consists of HfO2 with isolated inclusions of the molybdenum oxide. The inner layer contains oxycar(cid:2) bides of hafnium and silicon and grains of Mo5Si3C (in this region, no pores are present). Similar features were observed for the H5C(cid:2)II sample tested at 2400°C, although in this case the oxide layer (~300 μm thick) par(cid:2) tially splits off from the bulk part of the material. In spite of the presence of MoSi2 as the SiO2(cid:2)forming phase, no continuous glassy layer was found on the surface of HfC. The authors assume that it is precisely the pres(cid:2) ence of the matrix of carbide that prevents the forma(cid:2) tion of a stable layer of SiO2. It can be assumed that the molecules of CO that are formed upon the oxidation of HfC react with MoSi2 with the formation of volatile SiO. In the transverse section, the layered structure of the oxidized part is similar to that presented in the lit(cid:2) erature on the oxidation of monolithic HfC; the addi(cid:2) tion of MoSi2 to the composite [1] led to the formation  of SiO2, which partially filled the internal pores of the HfO2 layer, thus improving the oxidation resistance of the composite. In the inner layer, there are present particles of HfOxCy and SiOxCy; the latter are located between MoSi2 and HfC.  The experimental results for the HfC(cid:2)based sample suggest its partially catalytic behavior. At a tempera(cid:2) ture of 1800°C, the HfC(cid:2)based sample reveals noncat(cid:2) alytic properties. With  increasing temperature, the catalytic efficiency increases to 2 × 10-3 (2400°C).  The authors of [26] have studied the dynamic oxi(cid:2) dation of the ZrB2-15 vol % SiC ceramics (obtained by the method of hot pressing) under the conditions of aerothermal heating with the use of a high(cid:2)enthalpy supersonic flow of an N2/O2 mixture in a plasma(cid:2)wind tunnel. From the sample obtained by the method of electroerosion machining, a hemispherical sample (with a radius of curvature of 5 mm, see Fig. 22a [30]) was cut from the sample obtained by hot pressing. The hemispherical sample was placed on a holder made of Al2O3 at a distance of 10 mm from the edge of the torch; the gas expenditure was 1 g/s; the static pressure in the chamber, ~240 Pa; H0 and Pmax were varied in the  ranges of 9.3-18.5 MJ/kg and 6.8-9.5 kPa, respectively. The enthalpy H0 in the course of the experiment increased from 9.3 to 18.5 MJ/kg; the total time of heating was ~160 s. The emissivity coefficient is ~0.6 (at 1 μm; H0 = 18.5 MJ/kg). The catalytic effi(cid:2) ciency of recombination ε is (3-5) × 10-3. The tem(cid:2) perature of the surface of the sample used for the tests under the conditions of aerothermal heating reached 2130 K; Q in the vicinity of the stagnation point is equal to 5-7.5 MW/m2.  The surface of the sample after testing is covered by a glasslike SiO2(cid:2)based layer (the upper, very thin sub(cid:2) layer) which is obtained due to the oxidation of SiC, and by a layer of zirconium oxide (the lower sublayer, glued by a residual glasslike melt) (Fig. 23). The thick(cid:2) ness of the oxidized layer varies from 170-200 μm (near the stagnation point) to approximately 100 μm (on the lateral surface). Although the oxidized layer contains a certain amount of residual voids, it pre(cid:2)  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013      \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1685  (а)  R = 5 mm  (b)  R = 0.5 mm  (а)  sp  (b)  d1  d2  Blunt hemisphere  Sharp cone  (c)  (d)  Fig. 22. Photographs of model samples (a, b) before and (c, d) after tests in an electric(cid:2)arc plasma installation (R = radius of curvature) [30].  serves adhesion to the base layer and is sufficiently compact. An estimated partial pressure of O and O2 over the surface of the sample is 2 and 0.4 kPa, respec(cid:2) tively, which leads to the active oxidation of SiC. In this case, the oxidation of SiC occurs so rapidly that the sublayer depleted of SiC simply merely has no time to be formed and passes directly into the ZrO2 sub(cid:2) layer. The presence of small impurities of carbon is noted in the material.  The authors of [30] investigated the oxidation resis(cid:2) tance and the optical properties of the hot(cid:2)pressed ZrB2-SiC composite under the conditions of aero(cid:2) thermal heating in a flow of strongly dissociated air that simulates the conditions under which the appara(cid:2) tus is during the re(cid:2)entry. Samples of UHTCs with a sharp or blunt profile were subjected to the action of high(cid:2)enthalpy flows of an N2/O2 gas mixture (up to 10 MJ/kg) for 540 s; the surface temperature reached ~2100 K.  A ZrB2-20 vol % SiC composite was prepared by hot pressing from ZrB2 and β(cid:2)SiC using Si3N4 (3 vol %)  Table 4. Main parameters of arc(cid:2)jet plasma tests [30]  Fig. 23. Transverse sections of a sample after tests in an electric(cid:2)arc  plasma  installation  (SEM micrographs): (a) near the stagnation point, d1 ~ 200 µm; (b) far from the stagnation point, d2 ~ 160 µm [26].  as the sintering additive. From the thus(cid:2)obtained tab(cid:2) let, samples for the investigation of the aero(cid:2)thermo(cid:2) dynamic properties were cut, namely, a hemispherical sample (radius of curvature 5 mm) and a conical sam(cid:2) ple (radius of curvature 0.5 mm) (Fig. 22).  The model samples were subjected to the action of steady(cid:2)state(cid:2)enthalpy  flows using an arc(cid:2)jet plasma installation; the gas mixture consisted of N2 (80 wt %) and O2 (20 wt %); the mass flow rate of the mixture was 1 g/s; the distance from the model samples to the edge of the nozzle, 10 mm.  The experiments in the arc(cid:2)jet plasma installation were performed at a static pressure in the chamber equal to ~200 Pa; H0 and Pmax were varied in the ranges of 4.5-10.3 MJ/kg and 6.8-9.5 kPa, respectively (Table 4). The calculated values of the heat flux were 4.7 MW/m2  for  the  hemispherical  sample  and 11 MW/m2 for the conical sample. The samples were placed into a hot gas jet with H0 < 5 MJ/kg; then, H0 increased gradually to a maximum value of 10.3 MJ/kg; the total time of testing was ~10 min. The values of the temperature measured on the surface using an optical pyrometer (stage 5) were 2053 K (for the hemispheri(cid:2) cal sample) and 2083 K (for the conical sample); the coefficient of emissivity (at 1 μm) lay in the range of 0.6-0.65 at temperatures exceeding 2000 K.  After tests, the surface of the samples is covered by an oxidized layer, which consists of several sublayers (an upper SiO2(cid:2)based glasslike sublayer and a lower ZrO2/SiO2(cid:2)based sublayer). The thickness of the oxi(cid:2) dized layer formed in 230 s is from 60 to 150 μm for the hemispherical sample (Fig. 24a) and from 50 to 190 μm for the conical sample (Fig. 24b). In contrast to the con(cid:2) ical sample, for which the thickness of the SiC(cid:2)depleted region is maximum 70 μm at the apex of the cone, no  Arc power, kW  H0, MJ/kg  Pmax, kPa  τ, s  Stage 1  15.9  4.5  6.8/0.07 atm  80  Stage 2  Stage 3  Stage 4  18.9  5.5  7.4  60  22.4  7.3  8.1  60  25.6  8.6  8.8  60  Stage 5  29  10.3  9.5/0.1 atm  230  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1686  SIMONENKO et al.  Resin  Resin  (а)  1  2  Un(cid:3)oxidized balk  20 μm  1  2  3  4  (b)  25 μm  Fig. 24. Transverse sections of (a) the upper part of a hemispherical sample and (b) of the tip of a conical sample after testing (SEM): (1) upper glasslike layer with inclusions of particles of zirconium oxide of very small size; (2) oxide sublayer with coarse crystals of zirconium oxide; (3) region depleted of SiC; and (4) the unoxidized initial composition [30].  presence of a ZrB2 region depleted of SiC is revealed in the case of the hemispherical sample.  3.2.3. Tests  in which the model samples have a wedge shape with a radius of curvature less than 0.5 mm  [7, 11, 16, 17, 40]. The authors of [16] performed an interesting study on the action of the flame of an oxy(cid:2) acetylene torch on hot(cid:2)pressed wedge(cid:2)shaped samples of ZrB2-20 vol % SiC with sharp edges (Fig. 25) with different radii of curvature (R1, 0.15; R2, 0.5; R3, 1.0; R4, 1.5 mm). It has been established that, in accor(cid:2) dance with the known formula q ∝ R-1/2 (where q is the heat  flux),  the maximum  temperature Tmax  that observed at the tip of the wedge changed with increas(cid:2) ing radius of curvature (see Table 5). It was noted in [16]  that,  in  the course of  the experiment, Tmax decreased (especially strongly for the models R1 and R2, for which the most intensive loss of mass and abla(cid:2) tion of the material from the sharp tip was also noted, see Fig. 26). Based on the fact that, for the samples with a radius of curvature 0.15 and 0.5 mm, a sharp increase in this radius is observed as a result of the  (а)  Hot(cid:3)pressing direction  R  (b)  m  0 m 1  1 0 m mR1 = 0.15 mm  (c)  R2 = 0.5 mm  m  0 m 2  (d)  R3 = 1.0 mm  (e)  R4 = 1.5 mm  plasma(cid:2)chemical treatment, which leads to a change in  the calculated aerodynamic characteristics,  the radius R3 (1 mm) is recommended in this work as the optimum  geometry,  since  it  only  insignificantly changed during the experiment. The SEM results made it possible to conclude on the formation of a three(cid:2)layered oxidized zone: the first layer consists predominantly of ZrO2 with a small content of SiO2 (thickness 200-250 μm in the region of the maximum action); further, there is formed a SiC(cid:2)depleted zone; under this zone, the unaltered ZrB2-20 vol % SiC material is retained (Fig. 27).  The behavior of samples with a sharp leading edge with the radius of curvature 0.1 mm in the supersonic flow of plasma (electric arc plasmatron) for the hot(cid:2) pressed materials of substantially different nature was also investigated by the authors of [17]; they used (1) ZrB2-15 vol % SiC-2 vol % Si3N4 and (2) Si3N4- 35 vol % MoSi2-2 vol % Al2O3-5 vol % Y2O3. The second material was selected on the basis of the known excellent mechanical properties of the materials on the basis of Si3N4 (strength, cracks resistance); the measured values of the thermal conductivity for these materials differed considerably: these were 80 (25°C) and 58 (1500°C) W/(m K) for the first material, and  Fig. 25. Models with a sharp leading edge for ZrB2-SiC samples with different radii of curvature (R): (a) prepara(cid:2) tion of a model with a sharp leading edge from the initial sample (schematic); (b) R1, 0.15 mm; (c) R2, 0.5 mm; (d) R3, 1.0 mm; and (e) R4, 1.5 mm [16].  R1  R2  R3  R4  Fig. 26. The external appearance of samples after tests using an oxyacetylene torch [16].  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1687  31 (25°C) and 13 (1300°C) W/(m K) for the second material.  The nominal Mach number was equal to 3; the mass rate of the gas flow was 1 g/s; the static pressure in the chamber was ~200 Pa; the enthalpy varied from 8 to 16.5 MJ/kg; the power was increased gradually in steps. The ZrB2(cid:2)based sample was subjected to a two(cid:2) fold action with an increase in the mean temperature (without allowance for its extreme increase at the edge) to 1550 ± 25°C (the emissivity at 1 μm was 0.75-0.73). In the case of the Si3N4-MoSi2 sample, a catastrophic destruction occurred at the last step of the first cycle (Fig. 28); the authors of [17] explain this by insufficiently high thermal conductivity of the mate(cid:2) rial, which leads to a local overheating, softening of the material, and its ablation.  As can be seen, the white oxidized layer in the ZrB2(cid:2)based sample is formed only near the edge of the sample. The SEM study of the lateral surfaces (Fig. 29a) showed that the composition of the oxidized layer changes with approaching the sharp edge; at a distance of about 2 mm from the edge, ZrO2 is predominant on the surface; the subsequent zone (of about 1 mm wide) consists of island(cid:2)type structures containing ZrO2 dis(cid:2) tributed in the SiO2(cid:2)based glassy layer. At the very tip of the wedge, there is formed a porous layer of ZrO2  Zr  (f)  b  c  d  Int  C  O Zr  Zr  C1  Zr  Int  O C  R  Int  O C  C1  Table 5. Conditions of the plasma(cid:2)chemical treatment of samples with different radii of curvature [16]  Model Magnitude  of R, mm  Time  of ablation, s  R1  R2  R3  R4  0.15  0.5  1.0  1.5  300  300  300  300  Temperature  of the surface, °C  Tmax  2100  2040  2000  1930  Tavg  1935  1920  1910  1910  approximately 100-140 μm thick, which exhibits very low adhesion to the inner layers. The study of the microstructure of a polished section made it possible to conclude that the thickness of this ZrO2 layer varies from 140 to 50 μm (on the lateral surface) (Fig. 29b). The radius of curvature increased from 0.1 to 0.14 μm as a result of the plasma(cid:2)chemical treatment. The for(cid:2) mation of a SiC(cid:2)depleted region  is also observed, which has a significant thickness only on the side of the sharp edge. Thus, it has been shown in [17] that it is important to use UHTCs with a high thermal con(cid:2) ductivity, which allows, due to conductive heat trans(cid:2) fer inside the samples of complex shape, to redistribute heat into less heated regions and to ensure heat emis(cid:2) sion into the environment, thereby partially removing heat from the regions of local overheating.  One more interesting work [40] concerns a study of the interaction of ZrB2-5 vol % Si3N4-20 vol % SiCf (short fibers of Hi(cid:2)Nicalon) material (prepared by the method of hot pressing and machined to form a sharp wedge with the radius of curvature 0.1 mm) with the flame of the torch of the arc(cid:2)jet plasmatron on the base of inert gases (mass flow rate of 1 g/s). To simulate the required aerodynamic action characteristic of air, oxy(cid:2)  a  (а)  (b)  (d)  100 μm  (c)  100 μm  100 μm  (e)  100 μm  (а)  (b)  Fig. 27. Microstructure of sample R3 with the radius of curvature 1.0 mm after treatment using an oxyacetylene torch [16].  Fig. 28. Samples of (a) Si3N4-35 vol % MoSi2-2 vol % Al2O3-5 vol % Y2O3 and (b) ZrB2-15 vol % SiC-2 vol % Si3N4 after the action of a supersonic flow of air [17].  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  \\x0c', '1688  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013  SIMONENKO et al.  gen was mixed with nitrogen in the mixing chamber after the torch. In the experiments, a supersonic nozzle with the nominal Mach number equal to 3 was used; the power was increased in steps to H0 = 8-16.4 MJ/kg (Table 6); the total time of treatment was ~17 min. The temperature along the lateral surfaces, as measured pyrometrically (from a spot 3 mm in diameter), was  found to fall sharply in time; the temperature of the sharp edge was simulated with the application of com(cid:2) putational hydrodynamics (Fig. 30).  Figure 31 [40] demonstrates the exterior appear(cid:2) ance of a model sample before and after testing; noticeable changes in the shape are absent, the profile was preserved; the end of the wedge became brighter,  (b)  (а)  E  D  G  E  (2450-1840)°C  (1840-1650)°C  (1650-1585)°C  Zone 1: zirconia, d1 ~ 2 mm ~  Edge  Resin  1  R*  d1  3  2  d2  Back(cid:3)end  4  Zone 2: silica + zirconia, d2 ~ 1 mm ~  0.1 mm  Fig. 29. (a) Microstructure (SEM) of the lateral surface with superimposed calculated temperatures, and the photograph (b, left) and the microstructure of the transverse section (b, right) of a ZrB2-15 vol % SiC-2 vol % Si3N4 sample after a twofold plasma(cid:2) chemical treatment: (1) ZrO2(cid:2)based zone; (2) unoxidized material; and (3) SiC(cid:2)depleted zone [17].  Table 6. Maximum value of enthalpy, time of each treatment, maximum achieved temperature, emissivity measured by a pyrometer, and the total time of the action of the flow for ZrB2-SiCf and ZrB2-SiCp [40]  Sample  Experiment  H0max, MJ/kg  Time, s  Tmax, °C  ε  Total time  ZrB2-SiCf  f1  13.8  285  1380  0.88  16 min 45 s  f2  17.0  330  1590  0.86  f3  12.3  120  1395  0.65  f4  17.0  270  1680  0.54  ZrB2-SiCp  p1  10  240  1780  0.63  4 min  1700  1600  1500  1400  1300  1200  1100  350  300  250  150 200 Time, s  0  100  50  1000  2600  2400  2200  2000  1800  1600  7  10  9  8  5 6 x, mm  0  3  1400  4  2  1  T  e  m  p  e  r  a  t  u  r  e  ,  °  C  f4  f2  f1  f3  Experimental data k = 2 k = 5 k = 10 k = 66  (a)  (b)  Fig. 30. (a) Temperature conditions used during tests and (b) the results of the simulation of the surface temperature depending on the distance to the edge [40] (k, different thermal conductivities in the near(cid:2)surface zone with allowance for oxidation with the formation of ZrO2). Vertical line indicates the distance at which the temperature was measured by a pyrometer.    \\x0c', 'PROMISING ULTRA(cid:2)HIGH(cid:2)TEMPERATURE CERAMIC MATERIALS  1689  C  m m  5 1  (а)  (c)  SiO2  500 μm  (b)  50 μm  1  C O  0  1  2  5 μm  Si  1  Au  Au  2  C  O  Au  2  Si  Au  0  1  2  Fig. 31. The external appearance of the sample before and after plasma(cid:2)chemical treatment [40].  which, as is confirmed by the SEM data, indicates that ZrO2 is formed on the surface as the basic component.  A study of the microstructure of the initial sample showed that the Hi(cid:2)Nicalon fibers, which consist of nanosized SiC grains, amorphized Si-C-O phase, and residual carbon, were distributed in the composite material uniformly; as a result of hot pressing, they became structured into a SiC core surrounded by an oxycarbide SixCyOz shell, into which ZrC particles were built(cid:2)in. Between  the matrix grains,  second phases were formed, namely, ZrO2, BN, Zr-Si phase, and a glass, which contains Zr-Si-B-N-O [40].  Figure 32 displays the section of a sample in which at least three typical regions can be distinguished: (1) an external cracked layer of ZrO2 with inclusions of carbon agglomerates; (2) a ZrO2 layer (approximately 120 μm thick), which contains molten SiO2 in the regions, where SiC fibers were initially located (which was not observed when SiC particles rather than fibers were used); and (3) a region of ZrB2 depleted of SiC, with a thickness of ~400 μm. Note that the micro(cid:2) scopic ZrO2 particles are rounded rather than have a typical columnar morphology. The thickness of the external ZrO2 layer varied from 120 μm at the tip to 40-50 μm on the lateral surfaces [42]. The authors of [42] assume that the oxidation of multilayer fibrous formations in the composite can also occur according to reactions with the formation of carbon, especially, in the internal part, where the material consists of pure SiC and where a low oxygen pressure exists (reaction (5) and (6)). The oxidation of fibers in the zone under the ZrO2(cid:2)based layer begins, according to the authors of [40], from the passage of the external SixCyOz layer into the gas phase (with the formation of SiO, CO, and CO2); then, the oxidation of the SiC core with the for(cid:2) mation of SiO2 and gaseous products occurs, as in the  Fig. 32. Microstructure of the transverse section of the ZrB2-SiCf composite after four treatments in the plasma(cid:2) tron: (a) the external appearance and the scheme of distri(cid:2) bution of the layers upon going from the initial UHTC through the SiC(cid:2)depleted layer to the external ZrO2(cid:2)based layer containing both SiO2 and traces of carbon embedded into ZrO2; (b) an enlarged image of the end region with the precipitates of traces of C and SiO2 in ZrO2; and (c) the morphology of a SiC fiber below the depletion zone and the corresponding EDS spectra [40].  case of composites that contain fibers rather than par(cid:2) ticles of SiC. It is noted that, upon the oxidation of composites reinforced by SiCf, the fibers behave as single particles; they are oxidized more slowly than the globular SiCp particles that form a kind of a three(cid:2) dimensional carcass. The authors of [40] also believe that the formation of a noticeable amount of impurity carbon in the bulk of ZrO2 can lead to an increase in the thermal conductivity of the oxidized part of the material, which will positively affect the heat removal from the critical zones, decrease the overheating, soft(cid:2) ening, and, in the final analysis, ablation from the sharp tip.  CONCLUSIONS  In this work, we considered some aspects of heat transfer under  the  conditions of  the  interaction between components with a sharp leading edge and high(cid:2)enthalpy high(cid:2)speed flows of dissociated air; the results obtained make it possible to separate some characteristics of the materials, which would make them promising for the use in hypersonic flight vehi(cid:2) cles: (1) besides the high melting points and phase sta(cid:2) bility in a wide temperature range, the materials must have comparatively high oxidation resistance, in par(cid:2) ticular, in the reactions with atomic oxygen; 2) the material should have a high thermal conduc(cid:2) tivity, which must ensure heat removal from strongly  RUSSIAN JOURNAL OF INORGANIC CHEMISTRY   Vol. 58   No. 14   2013    \\x0c', '1690  SIMONENKO et al.  overheated regions; this is especially important for the samples with a sharp leading edge, since the local over(cid:2) heatings to ~2500°C cause softening and ablation by the high(cid:2)speed gas flows and lead to changes in the geometry of components, which cannot but affect the flight performance of articles on the whole; (3) the materials must have minimum catalytic activity γ in the exothermic reactions of the surface recombination  in order to minimize the chemical component of the aerodynamic heating; (4) the high emissivity coefficient makes it possible to intensify the process of the re(cid:2)radiation of heat obtained  from  the  relatively cold  lateral surfaces, which, along with  the high  thermal conductivity, makes possible for the material to play the role of a passive heat pipe, which provides the transfer of energy inside the system, and, as a result, from the system [13]. Specific features of the oxidation of materials based on ZrB2 and HfB2 alloyed by SiC have been analyzed; this composition of the system allows, due to the appearance of oxidation products such as Zr(Hf)O2, B2O3, SiO2,  the  formation of complex multilayer structures, where the external region (at comparatively low temperatures of 1650-1800°C) consists of a car(cid:2) cass of ZrO2 or HfO2, in the channels and on the sur(cid:2) face of which a viscous borosilicate glass is formed, which serves as a diffusion barrier for oxygen; and in the deeper  regions,  there can exist a  region of Zr(Hf)B2 depleted of SiC because of its active oxida(cid:2) tion. There have been considered methods (proposed by different research groups) of increasing the effi(cid:2) ciency of antioxidant layers both due to the introduc(cid:2) tion of additions (TaSi2, CrB2, etc.), which enhance the viscosity of borosilicate glasses and, consequently, decrease the diffusion through them, and, also, due to the introduction of some compounds (LaB6, La2O3), which must stabilize the arising Zr(Hf)O2, maximally level(cid:2)off the influence of polymorphic transforma(cid:2) tions upon repeated thermal loading, and decrease spalling in the incident gas flows. Finally, there have been described some works con(cid:2) cerning the behavior of samples under the action of high(cid:2)enthalpy flows of dissociated air, in particular, of samples with the shape that simulates sharp leading parts and wing edges of hypersonic flight vehicles. It has been shown that the decrease in the radius of cur(cid:2) vature essentially  increases the temperature  in the region of critical braking, which, in the case of insuffi(cid:2) ciently high thermal conductivity, can  lead to the destruction of the model.   ACKNOWLEDGMENTS  This study was supported in part by the Russian Foundation for Basic Research (grants no. 12(cid:2)03(cid:2) 33005(cid:2)mol_a_ved and 13(cid:2)03(cid:2)12206(cid:2)ofi_m) by a grant of the President of the Russian Federation for Young Scientists (grant no. MK(cid:2)1435(cid:2)2013.3).  REFERENCES  1. R. Savino, M. De Stefano Fumo, L. Silvestroni, and D. Sciti, J. Eur. Ceram. Soc. 28, 1899 (2008).  2. M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, J. Mater. Sci. 39, 5887 (2004).  3. S. Tang, J. Deng, S. Wang, et al., Mater. Sci. Eng., A 465, 1 (2007).  4. L. Li, Y. Wang, L. Cheng, et al., Ceram. Int. 37, 891 (2011).  5. J. Li, S. Meng, J. Han, et al., Key Eng. Mater. 368- 372, 1761 (2008).  6. M. Gasch, D. 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},{
  "_id": 227,
  "PDF": "PROTECTIVE OXIDE LAYERS FORMED DURING ELECTROCHEMICAL OXIDATION OF HAFNIUM CARBIDE.pdf",
  "Text": "['Powder Metallurgy and Metal Ceramics, Vol. 48, Nos. 9-10, 2009   PROTECTIVE OXIDE LAYERS FORMED   DURING ELECTROCHEMICAL OXIDATION   OF HAFNIUM CARBIDE   V. A. Lavrenko,1 V. N. Talash,1,3   M. Desmaison-Brut,2 and Yu. B. Rudenko1   UDC 546.241:620.193   The composition of oxide nanofilm formed on the HfC anode under electrolysis of a 3% NaCl   solution at potentials between - 0.20 and +1.90 V is studied using the potentiodynamic method of   polarization curves and Auger electron spectroscopy and SEM methods. The film consists of the  upper layer formed in polymolecular chemosorption of O2 and Cl2 gases followed by the HfO2 + C  (1 : 1) layer. Two inner layers represent a 7 at.% O2 solid solution in HfC and a mixture of  HfC0.7O0.3 and HfO (7 : 1).  The oxide film 30-40 nm thick formed at potentials up to 1.35 V is   protective and ensures very high corrosion resistance of HfC without any further polarization.   Keywords: hafnium carbide, chemosorption, anodic oxidation, formation of protective nanofilm,   hafnium oxycarbide, hafnium oxides.   Although hafnium carbide is rather expensive, it is widely used in modern engineering to produce space   rocket nozzles, structural components of gas-cycle nuclear jet engines, high-performance thermoionic generators,   parts of power ion engines, and protective shields and control rods for nuclear reactors. Hafnium carbide is   sometimes oxidized, oxide films being formed on its surface.     We should examine the nature of the electrochemical oxidation of hafnium carbide at small potentials of   anodic polarization to model the initial stages of forming such films not only in contact with marine water splashes   but also with atmospheric oxygen.    The HfC electrochemical oxidation is still to be examined, while HfC high-temperature oxidation is the   subject of many papers. In particular, it was hypothesized in [1] that hafnium carbide and oxycarbide oxidized to  HfO2 through an intermediate phase of HfO2C (this hypothesis was not verified further). Two-stage HfC oxidation  in air to 1200°C was established in [2]. It was shown that oxycarbide HfCxOy  formed in the first stage and oxide   HfO2 in the second.    Bargeron and his coauthors [3, 4] conducted more detailed research into the oxidation of HfC between   1400 and 2060°C [3, 4] to reveal three oxide layers in the cross-section of oxidized samples. The lower layer  represents carbide HfC with oxygen being dissolved in its lattice, the upper, quite porous layer contains HfO2, and   the dense intermediate layer consists of hafnium oxide and carbon. It was shown that it was the intermediate layer   that served as a barrier to the diffusion of oxygen inside the sample.     1Frantsevich Institute for Problems of Materials Science, National Academy of Sciences of Ukraine, Kiev,   Ukraine. 2The University of Limoges, France.    3To whom correspondence should be addressed; e-mail: shtal@ipms.kiev.ua.   Translated from Poroshkovaya Metallurgiya, Vol. 48, No. 9-10 (469), pp. 133-139, 2009. Original article   submitted January 22, 2009.   1068-1302/09/0910-0595 ©2009 Springer Science+Business Media, Inc.   595  \\x0c', 'Intensity  6000  4000  2000  0 10  20  30  40  50  60  70  2  , deg  Fig. 1. Diffraction pattern of the starting HfC sample   It was further established [5] that the starting carbon material could be reinforced during the chemical   deposition of hafnium carbide above 3000°C, forming a highly porous ‘HfC-carbon foam’ coating (open porosity   P = 98%). This seems to have served as the basis for developing the patented material for a nonconsumable   electrode for plasma arc welding [6], which represents a layer of hafnium oxycarbide under a graphite layer.     The group led by Shimada [7, 8] studied the high-temperature oxidation of HfC [7, 8] and analyzed   different planes of grown hafnium carbide single crystals. In particular, it is established in [7] that carbon forms   when HfC single crystals are oxidized at quite high temperatures and extremely low partial pressures of oxygen.   The paper [8] used thermal gravimetry (TG), differential thermal analysis (DTA), and mass spectroscopy to  establish that oxidation proceeded in two stages with two exothermal peaks at all PO2 considered; 100% oxidation,   i.e., a higher-temperature DTA peak, is characterized by carbon release.    However, the most interesting results on the interaction mechanism are reported in [9, 10]. For example, the   oxidation of HfC plane (100) was examined in [9] with photoemission spectroscopy using synchrotron radiation. It   was established that holding plane (100) in pure oxygen at room temperature causes it to interact with carbon atoms   and form CO molecules. Only after CO desorption do Hf atoms react with oxygen to form more than 10 loose  surface  layers containing HfO2. The oxidation of  the same plane was examined  in  [10] at 1100°C with  photoemission spectroscopy. The minimum electronic work function is associated with the formation of HfCxOy on   some part of the sample, while the other part of its surface is covered with a layer of HfO.   Shimada’s group obtained a two-layer oxide film during 4 h oxidation of HfC plane (200) at 700 to 1500°C   and oxygen pressure of 0.08 to 80 kPa [10]. The kinetics of its formation was examined with an electronic   microbalance, the surface structure with electron microscopy, and quantitative elemental composition of the layers   with a long-wave x-ray dispersion microanalysis.  It was established that the scale formed on the plane consisted of two zones. Zone 1 belonging to HfС  sample is porousless black compact scale containing, in addition to hafnium oxide, 25 at.% unoxidized carbon,  while external, very porous zone 2 represents monocline HfO2. Zone 1 is separated from the sample base with a   carbon-containing layer identified with x-ray spectroscopy and transmission electron microscopy.   This paper examines the kinetics and mechanism of anodic oxidation of HfC samples in a 3% NaCl   solution imitating marine water. The samples were produced by high-temperature isostatic pressing at the Chemical   Department at the University of Limoges. Figure 1 shows diffraction pattern of the starting HfC sample taken using   a DRON-2M diffractometer. It is seen that the sample contains admixtures whose content is more than 5%.   The oxidation kinetics was   examined using   a PI-50-1 potentiostate   and   anodic potentiodynamic   polarization curves with potential frequency of 5 mV/sec. Auger electron spectroscopy (LAS-2000 Riber) was used   for layer-by-layer identification of the composition of oxide films. The composition of each layer of the HfC film   formed in anodic oxidation (at.%) was determined by successive etching of the surface with Ar+ ions; etching   duration is 10 min. The content of elements in each layer was calculated using characteristic concentration peaks at   the following energies of Auger electrons E, eV: 150 (Hf), 190 (C,) 500 (O), 310 (N), 1080 (Cl).    Figure 2 shows the anodic polarization curve of the oxidation of the HfC sample after preliminary   mechanical treatment of its surface, washing with ethanol, and drying.  The oxidation curve of titanium carbide is    596    θ   \\x0c', 'HfC  TiC  E, V  0 2.  0 3.  0 8.  1 3.  1 8.  Fig. 2. Anodic polarization curves of the oxidation of the HfC sample and of the TiC sample (for   lg( )i  8  6  4  2  0  reference)   shown for reference for the same electrolysis conditions. All potentials are given with regard to the standard silver chloride reference electrode. Note higher corrosion resistance of the HfC sample in a 3% NaCl solution as   compared with the quite stable TiC sample (about one order of magnitude higher).   Our interpretation of the mechanisms in different stages of HfC electrochemical oxidation is based on   quantitative Auger electron spectroscopy (Table 1).    Figure 2 shows that the desorption of chemisorbed oxygen from the sample is observed at anodic potentials   between -0.20 V and +0.55 V; after which a dense protective oxide nanofilm about 40 nm thick, which consists of a  mixture of HfC0.7O0.3 and HfO, starts forming at 0.65 V. The formation of the film is completed at anodic potential   of 1.30 V. This potential value is optimal for obtaining a stable oxide film on the HfC samples to prevent their   further oxidation in air, possibly, even at quite high temperatures.    This stage of HfC oxidation can be expressed by the following equation:   8HfC + 12H2O = HfO + 7HfC0.7O0.3 + 3CO3 2- + 24H+ + 18e-.   (1)   TABLE 1. Content of Elements in Different Layers of the Electrochemical Coating on HfC  after Electrolysis in a 3% NaCl Solution   Distance from the layer   to the coating, nm    Surface    0.5    1.0    2.0    3.0    4.0    5.0    6.5    8.0  10.0  12.5  15.0  20.0  30.0  50.0   Content of elements, at.%    C   5.4  11.3  14.4  20.3  26.6  26.0  30.0  30.4  30.3  30.1  30.8  30.8  30.2  30.2  46.9   O   70.4  70.8  69.5  59.6  48.2  48.0  20.0  19.8  19.9  19.9  19.5  19.7  20.0  19.5   7.4   N   0.5  1.3  0.5  0.3  -  -  -  -  -  -  -  -  -  -  -   Hf     4.9  12.2  14.7  19.7  25.1  26.0  50.0  49.8  49.8  50.0  49.6  49.5  49.7  50.3  45.7   Cl   18.8    4.9    0.9    0.1  -  -  -  -  -  -  -  -  -  -  -   597    − − − − −     \\x0c', 'HfO  HfC  0.7  O  0.3  C  HfO  Fig. 3. Microstructure of HfC surface oxidized at 1.30 (a) and 1.90 V (b)   200  mμ  10  10 мкм  mμ  Figure 3a shows an electron microscopic image of the HfC sample oxidized at this stage. It is seen that the   ratio between the amount of finer, dark HfO crystallites and light crystalline precipitates of hafnium oxycarbide is   about 1 : 7.   If the oxidation process is not terminated, Hf4+ ions go into the electrolyte solution at anodic potentials   between 1.30 to 1.48 V and HfO2+ ions between 1.48 and 1.60 V:    HfC + 3H2O = Hf4+ + CO3  2- + 6H+ + 8e-,   HfC + 4H2O = HfO2+ + CO3  2- + 8H+ + 8e-.   (2)   (3)   This is indicative of the partial loss of integrity of the oxide film on HfC.    Finally, there is a new layer of unstable porous coating on the sample at potentials between 1.60 and  1.90 V, which represents a mixture of HfO2 and carbon in the approximate ratio 1 : 1. Figure 3b shows rare and  quite large crystalline aggregates against the dense HfO + HfC0.7O0.3 film, where finer black graphite crystals and  gray HfO2 crystals are closely grown together. This stage of anodic oxidation can be described by the following   electrochemical equation   2HfC + 5H2O = HfO2 + Hf4+ + C + CO3 2- + 10H+ + 12e-.   (4)  However, according to Auger electron spectroscopy, the porous, almost foamy, very thin (to 2 nm) external   film actually causes the HfC sample to lose its extremely high corrosion resistance because of the deep polarization   of the electrolyte imitating marine water.    The layer-by-layer Auger spectroscopy of the oxide coating formed in different stages of HfC anodic   oxidation (Table 1) has completely confirmed the above-proposed mechanism of stage-by-stage scaling on HfC.   Therefore, four oxide layers form on the HfC sample after anodic oxidation. The first, upper layer to 2 nm   in thickness contains not only chemisorbed oxygen (to 71 at.%) but also chemisorbed nitrogen (to 1.3 at.%). The   equilibrium redox potential of anodic oxygen deposition by reaction    2H2O ↔ O2↑ + 4H+ + 4e-   (5)   is +1.23 V relative   to   the standard hydrogen electrode;   i.e., 1.45 V relative   to   the reference chlorine-silver   electrode. However, this process occurs only at E > 1.95 V in our case because of the overpotential of oxygen   deposition on HfC. The electrochemical deposition of gaseous chlorine is observed at the same potentials on the  surface:   2Cl- ↔ Cl2 + 2e-,   (6)   whose equilibrium potential in the electric series of metalloids is +1.36 V relative to the standard hydrogen   electrode [12]. Table 1 shows that the upper adsorbed layer can contain to 19% of gaseous chlorine.    The second layer of the film is quite thin (to 2 nm) and porous, with ‘spots’ occupying no more than one  eighth of the sample surface (Fig. 3b). It contains white HfO2 crystallites closely grown together and black carbon   crystallites in the ratio 1 : 1.   598      \\x0c', 'Only does the third layer 30-40 nm thick, which consists of a mixture of hafnium oxycarbide HfC0.7O0.3   and lower black oxide HfO, have protective properties. According to Auger electron spectroscopy, it contains   ~30 at.% carbon, ~50 at.% hafnium, and 20 at.% oxygen. Hence, an oxide film with excellent protective properties   can be obtained on the HfC anode in electrolysis of a 3% NaCl solution with increasing potential from -0.10 V to   1.30 V and termination of the process when it reaches its maximum.  The fourth layer ≤15 nm thick, which is bounding between the initial composition of the sample and of the   third layer, represents a solid solution of oxygen (~7 at.%) in HfC (Table 1).  The research has shown that HfC has as high resistance to electrochemical corrosion as HfB2 that we   examined previously [12].   REFERENCES   1.   2.   3.   4.   5.   6.  7.   8.   9.   10.   11.   12.   V. A. Zhilyaev, Yu. G. Zainulin, S. I. Alyamovskii, and G. P. Shveikin, “High-temperature oxidation of   zirconium and hafnium oxycarbides and oxycarbonitrides,” Powder Metall. Met. Ceram., 11, No. 8, 632-  636 (1972).     R. F. Voitovich and E. A. Pugach, “High-temperature oxidation of ZrC and HfC,” Powder Metal. Metal   Ceram., 12, No. 11, 916-921 (1973).   B. Bargeron, C. B. Benson, R. C. Newman, and W. Robby, “Oxidation mechanisms of hafnium carbide and   hafnium diboride in the temperature range 1400-2100°C,” John Hopkins APL Technical Digest, 14, No. 1,   29-36 (1993).   B. Bargeron, C. B. Benson, A. N. Jette, and T. E. Phillips, “Oxidation of hafnium carbide   in   the   temperature range 1400-2060°C,” J. Am. Ceram. Soc., 76, No. 4, 1040-1046 (1993).   P. Sourdiaucourt, A. Derre, P. Delhaes, and P. David, “Mechanical reinforcement of carbon foam by   hafnium carbide deposit,” J. Phys. IV France, 9, 1187-1194 (1999).   Non-consumable Electrode for Plasma-Arc Welding, US Patent 4304984.   S. Shimada, F. Janazar, and S. Otani, “Oxidation of HfC and TiC single crystals with formation of carbon   at high temperatures and low oxygen pressures,” J. Am. Ceram. Soc., 83, No. 4, 722-728 (2000).   S. Shimada, “Thermoanalytical study on the oxidation of ZrC and HfC powders with formation of carbon,”   Solid State Ionics, 149, No. 3-4, 319-326 (2002).   K. Edamoto, Y. Shiroton, T. Sato, and K. Ozama, “Photoemission spectroscopy study of the oxidation of   HfC (100),” Appl. Surf. Sci., 244, No. 1-4, 174-177 (2005).   S. Shimada, “Formation of oxide layer on HfC (100) surface studied by photoemission spectroscopy,” J.   Surf. Sci. Nanotechn., 4, 219-226 (2006).    N. A. Izgaryshev and S. V. Gorbachev, Course of Theoretical Electrochemistry [in Russian], Gostekhizdat,   Moscow (1951), p. 503.   V. A. Lavrenko, V. N. Talash, M. Desmaison-Brut, et al., “Kinetics and mechanism of the electrochemical   oxidation of hafnium boride,” Powder Metall. Met. Ceram., 48, No. 7-8, 462-465 (2009).   599        \\x0c']"
},{
  "_id": 228,
  "PDF": "Pursuing enhanced oxidation resistance of ZrB2 ceramics by SiC and WC co-doping.pdf",
  "Text": "['Journal of the European Ceramic Society 38 (2018) 5311-5318  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e l sev ie r .com / loca te / jeu rce ramsoc  Original Article  Pursuing enhanced oxidation resistance of ZrB2 ceramics by SiC and WC codoping  T  Xiaoqiang Fenga,b,c, Xin Wanga,b,c,⁎, Yuan Liua,b,c, Wei Tiand, Min Zhanga,b,c, Xian Jiand, Liangjun Yind, Linbo Zhanga,b,c, Jianliang Xiea,b,c, Longjiang Denga,b,c  a National Engineering Research Center of Electromagnetic Radiation Control Materials, University of Electronic Science and Technology of China, 2006 Xiyuan Road, Chengdu 611731, PR China b State Key Laboratory of Electronic Thin Film and Integrated Devices, University of Electronic Science and Technology of China, 2006 Xiyuan Road, Chengdu 611731, PR China c Key Laboratory of Multi-Spectral Absorbing Materials and Structures of Ministry of Education, University of Electronic Science and Technology of China, 2006 Xiyuan Road, Chengdu 611731, PR China d School of Materials and Energy, University of Electronic Science and Technology of China, 2006 Xiyuan Road, Chengdu 611731, PR China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Ultra-high temperature ceramics Zirconium diboride Co-doping Oxidation resistance  The oxidative degradation of ZrB2 ceramics is the main challenge for its extensive application under high temperature condition. Here, we report an effective method for co-doping suitable compounds into ZrB2 in order to significantly improve its anti-oxidation performance. The incorporation of SiC and WC into ZrB2 matrix is achieved using spark plasma sintering (SPS) at 1800 °C. The oxidation behavior of ZrB2-based ceramics is investigated in the temperature range of 1000 °C-1600 °C. The oxidation resistance of single SiC-doped ZrB2 ceramics is improved due to the formation of silica layer on the surface of the ceramics. As for the WC-doped ZrB2, a dense ZrO2 layer is formed which enhances the oxidation resistance. Notably, the SiC and WC co-doped ZrB2 ceramics with relative density of almost 100% exhibit the lowest oxidation weight gain in the process of oxidation treatment. Consequently, the co-doped ZrB2 ceramics have the highest oxidation resistance among all the samples.  1.  Introduction  Boride, carbide and nitride ceramics, such as TaC, ZrB2, ZrC, HfB2, HfC and HfN, are called as ultra-high temperature ceramics (UHTCs) owing to their high melting temperatures up to 3000 °C [1,2]. Due to their unique characteristics of high melting point, hardness, good thermal shock resistance and excellent oxidation resistance, the UHTCs are considered as promising materials for application in thermal protective systems (TPS) [3-6]. Among the UHTC materials, ZrB2 has attracted considerable attention due to its many advantages, including (6.09 g/cm3), low theoretical density high electrical conductivity (10.3 × 104 S/cm) and thermal conductivity (58.2 W/(m K)−1) [7-10]. However, ZrB2 starts to undergo oxidation at about 700 °C (Reaction (1)) and is severely degraded at temperatures above 1100 °C. At temperatures below 1100 °C, molten B2O3 covers the surface of ZrB2 to inhibit further oxidation. However, the anti-oxidation protection fails at 1100 °C due to the rapid evaporation of B2O3 gas (Reaction (2)) and  therefore, only a porous and non-protective ZrO2 layer remains on the surface [7].  ZrB ( c  2  )  +  5 2  O ( g  2  )  ZrO ( c  2  )  =  +  B O ( ) l  2  3  B O ( ) =l  2  3  B O ( g  2  3  )  (1)  (2)  Several methods have been reported to improve the oxidation resistance of ZrB2, and all of these strategies aim to form an anti-oxidation protective layer on the surface of ZrB2. These methods can be divided into three groups: (1) forming a continuous glass layer; (2) densifying the porous ZrO2 layer to inhibit oxygen diffusion; (3) combining the two effects to increase the energy barrier for oxygen diffusion. In the first method, researchers usually use silicon-containing additives (i.e. SiC, MoSi2 and Si3N4) [11-15], since SiO2 has higher viscosity, higher melting temperature and lower vapor pressure than B2O3 [15]. In the second method, WC is one of the typical dopants. The oxidation product WO3 is present as liquid phase in ZrO2 and promotes densification of  ⁎ Corresponding author at: National Engineering Research Center of Electromagnetic Radiation Control Materials, University of Electronic Science and Technology of China, 2006 Xiyuan Road, Chengdu 611731, PR China. E-mail address: xinwang@uestc.edu.cn (X. Wang).  https://doi.org/10.1016/j.jeurceramsoc.2018.07.041 Received 26 March 2018; Received in revised form 25 July 2018; Accepted 26 July 2018  Available online 27 July 2018 0955-2219/ © 2018 Elsevier Ltd. All rights reserved.  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  the porous ZrO2 layer via liquid phase sintering [16,17]. In the third method, both SiC and WC are used to improve the oxidation resistance of ZrB2 ceramics [5]. The oxidation mechanism of single-doped ZrB2 ceramics has been well studied, but the oxidation mechanism of codoped ZrB2 ceramics is still unclear due to the complex interaction among the various oxidation products. Herein, we investigate the co-doping approach to improve the oxidation resistance of ZrB2 ceramics. The morphology, composition and structural evolution of the oxidation products are studied carefully to clarify the oxidation mechanism of ZrB2-SiC-WC compounds. For comparison, the oxidation behaviors of pure ZrB2 ceramics and singledoped ZrB2 ceramics are also studied.  2. Experimental procedure  2.1. Sample preparation and characterization  Raw materials including ZrB2, SiC and WC are commercial chemi(10-15 μm, 99.5%) and SiC (1.6-1.7 μm, 99.5%) powders cals. ZrB2 were purchased from Eno Material, China. WC (2.9-3.1 μm, 99.5%) powders were obtained from Kelong Chemical Reagent Factory, China. The as-designed compositions of synthesized powders are given in Table 1. The powder mixtures are ball-milled in ethanol for 1 h. These powders are then dried for 24 h at 80 °C. ZrB2 ceramic composites are prepared by spark plasma sintering (SPS 4-6-20, Chen Hua) using a graphite mould at 1800 °C for 5 min under 40 MPa of uniaxial pressure in vacuum. The furnace is heated up to the sintering temperature at the heating rate of 100 °C/min. After holding the sintering temperature for 5 min, the furnace is naturally cooled down to room temperature. The sintered density of the specimens is determined using the Archimedes method with deionized water as the immersion medium. The theoretical density of each specimen is calculated based on the volume fractions of ZrB2, SiC and WC in the ceramics. Crystalline phases, microstructures and chemical compositions of ZrB2 ceramics are characterized by using X-ray diffraction (XRD, XRD-7000, Shimadzu), scanning electron microscopy (SEM, JSM-7600F, JEOL) and energy-dispersive spectroscopy (EDS, NORAN SYSTEM 7, Thermo Scientific), respectively.  2.2. Oxidation tests  Before the oxidation treatment, all the bulk specimens are polished using abrasive paper and are ultrasonically cleaned in ethanol afterwards. The oxidation process is conducted in a tabular furnace with molybdenum disilicide (MoSi2) as heating element. All ZrB2-based ceramics are oxidized at 1000 °C, 1200 °C, 1400 °C and 1600 °C in air. Subsequently, the oxidized ZrB2-based ceramics are embedded in epoxy resin and the cross-sections of ZrB2-based ceramics are polished by a polisher (LaboForce-100, Struers). Thermogravimetric analyses (TGA) of ZrB2-based ceramics are carried out using a thermal gravimetric  Table 1  List of as-designed ZrB2-based ceramics. ZrB2 (90 vol%)-SiC(10 vol%) composite is named as ZS10, similarly, ZrB2 (96.57 vol%)-WC(3.43 vol%) composite is named as ZW3. Other sample is defined in the similar way shown in the Table.  Sample  Pure ZrB2 ZS10 ZS20 ZS30 ZW3 ZW7 ZW11 ZS20W5  Composition  ZrB2(vol%)  100 90 80 70 96.57 93.02 89.35 75  SiC(vol%)  WC(vol%)  0 10 20 30 0 0 0 20  0 0 0 0 3.43 6.98 10.65 5  5312  analyzer (STA 449C, Netzsch). Specimens are heated at 10 °C/min up to 1400 °C in air atmosphere at a flow rate of 20 ml/min to obtain nonisothermal thermo gravimetric curve. Specimens are also heated at 40 °C/min up to 1200 °C and held for 2 h in air at the flow rate of 20 ml/ min to obtain isothermal thermo gravimetric curve.  3. Results and discussion  3.1. Microstructure of  the as-sintered ZrB2-based ceramics  The addition of SiC and WC has a significant influence on the microstructure of ZrB2-based ceramics. It is found that pure ZrB2 and the ZrB2 composites differ from each other in terms of their microstructure and density. The measured relative densities of all ZrB2 composites are higher than that of pure ZrB2, which is helpful for the oxidation resistance. Compared with the single(SiC or WC) doped samples, SiC and WC co-doped ZrB2 composites have the highest relative density of up to almost 100%, indicating the feasibility of co-doping approach using SPS technology. Compared with the normal pressureless sintering method, the density of pure ZrB2 is increased by 20% and the relative density value is 72% on average, which is still relatively low. Next, SEM observations are applied to confirm the microstructure of ZrB2-based ceramics and understand the density changes. As shown in Fig. 1-a, the pure ZrB2 bulk has a large number of porosities in the size range of 1 μm-9 μm resulting in relatively lower density. After SiC and WC doping, the porosity of ZrB2 composite decreases with the pore size range of 0.5 μm-3 μm, as shown in Fig. 1-b and Fig. 1-c, respectively. This change in porosity obviously favors the increase in relative density. From Fig. 1-b, it can be seen that SiC particles in darker color are fairly uniformly distributed in ZrB2 matrix and fill up the pore space between grain boundaries of ZrB2 particles, leading to decrease in ZrB2 porosity. Samples with higher amounts of doped SiC (10 vol%, 20 vol% and 30 vol%) have greater relative densities of 85%, 93% and 95%, respectively. Similarly, all the ZrB2-WC composites containing 3 vol %-11 vol% of WC display higher relative density. It is found that the optimal doping ratio is 6.98 vol% of WC (ZW7 sample), showing the highest value of 97%. The relative densities of ZW3 and ZW11 are 80% and 87%, respectively. The WC particles are also uniformly distributed in ZrB2 matrix as shown in Fig. 1-c. The ZrB2 particles are connected to each other closely. The reason is that WC plays a role in eliminating the oxygen contamination in ZrB2 matrix [18]. Thus, doping with WC favors the densification of the ZrB2 ceramics to improve the relative density. Importantly, after SiC and WC co-doping, the as-designed composite of ZrB2-20 vol%SiC-5 vol%WC (ZS20W5) has the relative density of almost 100% and the morphology is shown in Fig. 1-d. Both SiC and WC particles are uniformly distributed in ZrB2 matrix. Moreover, the pores between the grain boundaries and impurities in the ZrB2 matrix are eliminated due to the complementary effect of SiC and WC, which is helpful to improve the oxidation resistance. The XRD patterns of ZrB2-based ceramics shown in Fig. 2 confirm the hexagonal phase of ZrB2 (PDF #34-0432) as the matrix. The XRD pattern of the ZrB2-SiC composite of ZS10 is similar to that of pure ZrB2. The SiC peaks are not observed for the ZS10 sample because of the small amount of SiC. On the other hand, some SiC diffraction peaks are recorded for the ZrB2-SiC composites of ZS30 and ZS20W5, which confirms the existence of doped SiC. However, WB and ZrC are indexed after adding WC into the ZrB2 matrix. Interestingly, there is no WC phase in the as-prepared ZrB2-WC samples, which suggests that the added WC forms a solid solution with ZrB2 [16].  3.2. Microstructure evolution in the oxidation process  3.2.1. Evolution of oxidized surface The microstructure evolution of the oxidized surface of ZrB2-based ceramics is considered. As shown in Fig. 3, the surface of pure ZrB2 is  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  Fig. 1. Surface morphology of ZrB2-based ceramics: (a) pure ZrB2, (b) ZS30, (c) ZW11, (d) ZS20W5.  Fig. 4. For all oxidized samples, relatively dark and bright colored regions are found on the surface, which exhibit different morphologies and size distributions. The relatively dark substances are observed in Fig. 4-a-c due to the existence of SiO2 in ZrB2-SiC composites. After the oxidation of ZS10 at lower temperature of 1200 °C, the dark substances are small sized and distributed uniformly, as shown in Fig. 4-a. With the increase in oxidation temperature up to 1400 °C and 1600 °C, the dark regions become larger, as shown in Fig. 4-b and -c, respectively. It is confirmed that more silica is formed randomly on the surface of ZS10 at higher oxidation temperature. Furthermore, some gaseous bubbles are observed on the surface after oxidation at 1600 °C for high SiC-content (more than 20 vol%) doped ceramic, as shown in Fig. 4-d. The formation of the bubbles is related to the evolution of gas-phase products (i.e. CO, CO2) [19]. The elevated temperature causes an increase in oxidation rate and the rapid growth of silica. Therefore, the coherent silica layer exerts both positive and negative effects on the oxidation process. For the positive effect, the silica layer inhibits the oxygen diffusion to slow down the oxidation rate. However, the gas-phase products are also sealed under the silica layer. As the inner pressure becomes high enough and the gas leaks out, the silica layer is destroyed, which increases the risk of continuous oxidation of ZrB2 ceramic. Fig. 5 shows the XRD patterns of ZrB2-based ceramics after oxidation at 1000 °C-1600 °C for 3 h. It is found that the major phase is ZrO2 (PDF #37-1484) for these ZrB2-based ceramics. However, minor ZrB2 peaks still remain in ZS30 and ZS20W5 after oxidation at 1000 °C, which suggests that SiC has a significant effect on improving the oxidation resistance of ZrB2-based ceramics. Additionally, no WO3 is  Fig. 2. XRD patterns of ZrB2-based ceramics: (a) pure ZrB2, (b) ZS10, (c) ZS30, (d) ZW3, (e) ZW11, (f) ZS20W5.  oxidized completely. All surface particles are cracked into small fragments due to the rapid evaporation of B2O3 and volume change of ZrO2, indicating the poor anti-oxidation performance of pure ZrB2. Different treatment temperatures including 1200 °C, 1400 °C and 1600 °C are adopted to investigate the oxidation behavior of both ZS and ZSW samples. The typical results of ZS10 and ZS20W5 are shown in  Fig. 3. Surface SEM images of pure ZrB2 after oxidation (a) at 1000 °C for 3 h, (b) larger version of (a).  5313  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  Fig. 4. Surface SEM images of ZrB2-based composites after oxidation for 3 h: (a) ZS10 oxidation at 1200 °C, (b) ZS10 oxidation at 1400 °C, (c) ZS10 oxidation at 1600 °C, (d) ZS20W5 oxidation at 1600 °C.  detected in ZW11 ceramic, as shown in Fig. 5-c, which indicates that most of the WO3 on the surface is evaporated. ZrSiO4 phase appears in ZS20W5 after oxidation at 1400 °C for 3 h, which can improve the oxidation resistance of the ceramic [11]. The formation mechanism of ZrSiO4 is related to active oxidation of SiC and the reaction routes are  expressed in Eqs. (3) and (4) [20].  ZrO  2  SiO  2  ZrSiO  4  =  +  ZrO  2  SiO  +  +  1 2  O  2  ZrSiO  4  =  (3)  (4)  Fig. 5. XRD patterns of ZrB2-based ceramics after oxidation for 3 h at 1000 °C, 1200 °C, 1400 °C and 1600 °C: (a) pure ZrB2, (b) ZS20, (c) ZW11, (d) ZS20W5.  5314  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  Fig. 6. Cross-section images for ZrB2-based ceramics after oxidation for 3 h: (a) pure ZrB2 oxidation at 1000 °C, (b) ZS10 oxidation at 1400 °C, (c) ZS20 oxidation at 1400 °C, (d) ZS30 oxidation at 1400 °C.  3.2.2. Evolution of oxidation layer The cross-sectional backscattered electron images of oxidized ZrB2based ceramics are shown in Fig. 6. No layered structure is found in the oxidized pure ZrB2 and only a small amount of residual B2O3 is found on the surface. As seen from Fig. 6-b, the ZS10 sample is divided into three layers: (1) loose and porous ZrO2 layer with a small quantity of SiO2, (2) SiCdepleted layer, and (3) unoxidized ZrB2-SiC layer. The formation of the SiC-depleted layer is related to the active oxidation of SiC [9,15,21,22]. The thickness of oxide layer reaches up to 145 μm. A crack is present between the ZrO2 layer and SiC-depleted layer due to the phase transformation of ZrO2 and residual thermal stresses caused by the mismatch in coefficients of thermal expansion between ZrO2 and ZrB2 with the change in temperature. Fig. 6-c shows the cross-sectional micrographs of oxidized ZS20. Compared to ZS10, the oxidized layer structure of ZS20 shows the following differences: (1) the silica layer of ZS20 is thicker and reaches 26 μm; (2) a small amount of SiC is left in the second layer of ZS20, reasonably called SiC-poor layer, and the boundary crack is gone; and (3) the overall thickness of the oxidized layer (127 μm) is lower. Two oxide layers are observed in the cross-sectional micrographs of oxidized ZS30, as shown in Fig. 6-d. More silica is generated in ZrB2 doped with higher content of SiC, which makes the silica layer grow to 60 μm. This layer makes oxygen diffusion more difficult. Thus, the the oxidized layer is reduced to 110 μm. thickness of Fig. 7 shows the cross-sectional optical microscopy images of oxidized ZrB2-WC composites. Only a thin oxide layer is observed on the surface when the oxidation temperature is 1200 °C. Three different oxide layers are found on the surface of ZW3 after oxidation at 1400 °C for 3 h. The vapor pressures of WO3, (WO3)2, (WO3)3 and (WO3)4 at 1600 °C are 8.25 × 10−3 Pa, 75.2 Pa, 565 Pa and 20.4 Pa, respectively [16]. According to Reaction (5) [16,17,23], the rapid evaporation of WO3 in first layer (outermost layer) results in the porous structure of the oxidized ZW3 samples. In contrast, the second layer has a dense structure, due to the liquid phase sintering of WO3 with ZrO2, which acts as a barrier to oxygen diffusion [16]. The third layer is the WC-poor layer. However, the third layer disappears when the oxidation temperature rises to 1600 °C. Excessive doping of WC makes the oxide layer crack due to the huge volumetric expansion of WO3 during oxidation  and different thermal expansion coefficients of the oxidized surface and the unoxidized ceramic [17].  n  WO ( s  3  )  =  (WO )  3  n  ( g  )  (5)  From the cross-sectional backscattered electron image and elemental mappings of oxidized ZS20W5 shown in Fig. 8, two oxide layers are observed, similar to the oxidation behavior of ZS30. SiC and WC are oxidized gradually and the oxidation products move to the surface, as indicated by the Si and W mapping. The first layer is composed of SiO2, ZrO2 and WO3 and the second layer is SiC/WC-poor layer. The codoping of SiC and WC leads to a much denser SiC/WC-poor layer compared to the singly doped ZrB2 composites. The channels for oxygen transportation in SiC-depleted layer are reduced, and then oxygen diffusion is inhibited. Thus, oxidation resistance of the co-doped ZrB2 is improved. Furthermore, the cracks between the oxide layers almost disappear, which is beneficial to the oxidation resistance of ZrB2-based ceramics. The oxides thickness of ZrB2-based ceramics after oxidation at 1400 °C for 3 h is shown in Table 2. It can be seen that the compositions of ZrB2-based ceramics have an obvious effect on the thickness of oxides. The oxides thickness of ZrB2-based ceramics decreases with the increase in SiC content. The ZS20W5 has the lowest oxides thickness value of 102 μm, which shows that the ZS20W5 has the best performance of anti-oxidation among the ZrB2-based ceramics we investigated. Fig. 9 shows the diagram of oxidation layers of pure ZrB2, ZrB2-SiC composite, ZrB2-WC composite and ZrB2-SiC-WC composite. The doped SiC forms a protective layer of glass phase silica on the surface, while the doped WC forms a denser ZrO2 layer to inhibit the oxygen diffusion. Therefore, the co-doping of SiC and WC has a complementary effect on improving the oxidation resistance of ZrB2-SiC-WC composite.  3.3. Thermogravimetric analyses of ZrB2-based ceramics  The relationship between weight gain and oxidation time for ZrB2based ceramics during isothermal oxidation at 1200 °C is presented in Fig. 10-a-c. The curve shape is parabolic, as the oxidation behavior is controlled by the diffusion step and fits Jander model [24-26]. Table 3 presents the detailed analyses of the oxidized fraction of ZrB2-based  5315  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  Fig. 7. Cross-section images for ZrB2-WC composites after oxidation for 3 h: (a) ZW3 at 1200 °C, (b) ZW3 at 1400 °C, (c) ZW3 at 1600 °C, (d) ZW11 at 1400 °C.  ceramics, which is calculated from oxidation weight gain, as a function of volume content of SiC and WC. It can be seen that the oxidation resistance is significantly affected by the composition of ZrB2-based ceramics. The oxidation fraction of ZrB2-based ceramics decreases with the increase in SiC content due to the formation of more silica glass on the surface. An optimal content (6.98 vol%) of WC provides the best oxidation resistance of ZrB2-WC ceramics due to the dense oxide layer formed through the liquid phase sintering reaction of ZrO2 and WO3 [16]. Furthermore, the huge volumetric expansion results in great internal stress during the oxidation of WC, which extrudes the gas and eliminates some pores in ZrO2 layer. As a result, the diffusion rate of oxygen through ZrO2 layer is slowed down. Excessive doping (10.65 vol %) of WC deteriorates the oxidation performance of ZrB2-based ceramics, since the internal stress exceeds the limit and the protective oxide layer is destroyed. ZrB2 ceramics co-doped with SiC and WC (ZS20W5) have the lowest oxidized fraction value of 5%, indicating the best anti Table 2  Thickness of oxides for ZrB2-based ceramics.  Sample  Pure ZrB2 ZS10 ZS20 ZS30 ZW3 ZW7 ZW11 ZS20W5  Thickness of oxides (oxidized at 1400 °C for 3 h)  oxidized completely 145 μm 127 μm 110 μm 260 μm / 280 μm 102 μm  Fig. 8. Cross-section and elemental mapping for ZS20W5 after oxidation for 3 h: (a) oxidation at 1400 °C, (b) oxidation at 1600 °C, (c) O elemental mapping, (d) Si elemental mapping, (e) Zr elemental mapping, (f) W elemental mapping.  5316  \\x0c', 'X. Feng et al.  Journal of the European Ceramic Society 38 (2018) 5311-5318  Table 3  Oxidized fraction for ZrB2-based ceramics.  Sample  Pure ZrB2 ZS10 ZS20 ZS30 ZW3 ZW7 ZW11 ZS20W5  Isothermal oxidation  Non-isothermal oxidation  22% 11% 8% 6% 14% 8% 9% 5%  21% 7% 5% 2% 12% 3% 9% 1%  platform time corresponding to the slower weight gain stage, indicating the worst oxidation resistance. The sharp increase in weight gain at the final stage is due to the failure of the protective oxide layers. It is noted that the ZS20W5 composites have the lowest oxidized fraction of 1% during the non-isothermal oxidation, which indicates that ZS20W5 has the best oxidation resistance.  Fig. 9. The diagram of oxidation layers of four kinds of ZrB2-based ceramics: (a) pure ZrB2, (b) ZrB2-SiC composite, (c) ZrB2-WC composite and (d) ZrB2-SiC-WC composite.  4. Conclusion  oxidation performance. This means that SiC and WC together exert a complementary effect to improve the oxidation resistance of ZrB2-SiCWC composites. The weight gain with temperature for ZrB2-based ceramics during non-isothermal oxidation up to 1400 °C is given in Fig. 10-d-f. No weight change is observed for any of the ZrB2-based ceramics at lower temperatures ranging from 30 °C to 650 °C. The first mass change of ZrB2-based ceramics happens at around 650 °C, which is ascribed to the initial oxidation of ZrB2-based ceramics. After a relatively rapid weight gain during further heating, a slower weight gain stage occurs with temperature up to 1300 °C. Moreover, pure ZrB2 has the shortest  In this work, ZrB2-based ceramics are prepared by Spark Plasma Sintering at 1800 °C under argon atmosphere. The compositions of ZrB2-based ceramics have an obvious effect on the oxidation behavior. SiC improves the oxidation resistance of ZrB2 ceramics by forming a silica protective layer on the surface of the ceramics, while WC improves the oxidation resistance of ZrB2 ceramics by forming a dense internal oxidation layer. The co-addition of SiC and WC results in the highest density of almost 100%, the thinnest oxides thickness and the minimal oxidized fraction for ZrB2-20 vol%SiC-5 vol%WC composite, indicating the best oxidation resistance performance. The co-doped SiC and WC have a positive and complementary effect on the oxidation resistance of ZrB2-based ceramics due to the formation of protective SiO2 layer and the densification of ZrO2 layer.  Fig. 10. Mass increase of ZrB2-based ceramics, (a) isothermal oxidation of ZrB2-SiC under 1200 °C, (b) isothermal oxidation of ZrB2-WC under 1200 °C, (c) isothermal oxidation of ZrB2-SiC-WC under 1200 °C, (d) non-isothermal oxidation of ZrB2-SiC, (e) non-isothermal oxidation of ZrB2-WC, (f) non-isothermal oxidation of ZrB2SiC-WC.  5317  \\x0c', 'Journal of the European Ceramic Society 38 (2018) 5311-5318  [15]  [16]  [17]  [14]  [12] W.G. Fahrenholtz, The ZrB2 volatility diagram, J. Am. Ceram. Soc. 88 (2005) 3509-3512. [13] D. Sciti, M. Brach, A. Bellosi, Long-term oxidation behavior and mechanical strength degradation of a pressurelessly sintered ZrB2-MoSi2 ceramic, Scr. Mater. 53 (2005) 1297-1302. J. Han, P. Hu, X. Zhang, S. Meng, W. Han, Oxidation-resistant ZrB2-SiC composites at 2200°C, Compos. Sci. Technol. 68 (2008) 799-806. F. Monteverde, A. Bellosi, Oxidation of ZrB2-based ceramics in dry air, J. Electrochem. Soc. 150 (2003) B552-B559. S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Improved oxidation resistance of zirconium diboride by tungsten carbide additions, J. Am. Ceram. Soc. 91 (2008) 3530-3535. S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Oxidation of zirconium diboride with tungsten carbide additions, J. Am. Ceram. Soc. 94 (2011) 1198-1205. [18] H.-B. Ma, J. Zou, P. Lu, J.-T. Zhu, Z.-Q. Fu, F.-F. Xu, G.-J. Zhang, Oxygen contamination on the surface of ZrB2 powders and its removal, Scr. Mater. 127 (2017) 160-164. S. Gangireddy, S.N. Karlsdottir, S.J. Norton, J.C. Tucker, J.W. Halloran, In situ microscopy observation of liquid flow, zirconia growth, and CO bubble formation during high temperature oxidation of zirconium diboride-silicon carbide, J. Eur. Ceram. Soc. 30 (2010) 2365-2374. [20] D. Gao, Y. Zhang, J. Fu, C. Xu, Y. Song, X. Shi, Oxidation of zirconium diboride-silicon carbide ceramics under an oxygen partial pressure of 200Pa: formation of zircon, Corros. Sci. 52 (2010) 3297-3303. P.A. Williams, R. Sakidja, J.H. Perepezko, P. Ritt, Oxidation of ZrB2-SiC ultra-high temperature composites over a wide range of SiC content, J. Eur. Ceram. Soc. 32 (2012) 3875-3883. [22] W.G. Fahrenholtz, Thermodynamic analysis of ZrB2-SiC oxidation: SiC-depleted region, J. Am. Ceram. Soc. 90 (2007) 143-148. [23] K.N. Marushkin, A.S. Alikhanyan, J.H. Greenberg, V.B. Lazarev, V.A. Malyusov, O.N. Rozanova, B.T. Melekh, V.I. Gorgoraki, Sublimation thermodynamics of tungsten trioxide, J. Chem. Thermodyn. 17 (1985) 245-253. J.H. Sharp, G.W. Brindley, B.N.N. Achar, Numerical data for some commonly used solid state reaction equations, J. Am. Ceram. Soc. 49 (1966) 379-382. J.H. Sharp, S.A. Wentworth, Kinetic analysis of thermogravimetric data, Anal. Chem. 41 (1969) 2060-2062. [26] K.-C. Chou, X.-M. Hou, Kinetics of high-temperature oxidation of metallic materials, J. Am. Ceram. Soc. 92 (2009) 585-594.  inorganic non formation of a  [24]  [25]  [19]  [21]  X. Feng et al.  Acknowledgments  This work was supported by the NSFC (Grant Nos. 51702041) and the Open Foundation of Key Laboratory of Multi-spectral Absorbing Materials and Structures, Ministry of Education (ZYGX2016K009-3).  References  [5]  [3]  [1] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000°C + hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887-5904. [2] M.J. Gasch, D.T. Ellerby, S.M. Johnson, Ultra High temperature ceramic composites, in: N.P. Bansal (Ed.), Handbook of Ceramic Composites, Springer US, Boston, MA, 2005, pp. 197-224. S.-Q. Guo, J.-M. Yang, H. Tanaka, Y. Kagawa, Effect of thermal exposure on strength of ZrB2-based composites with nano-sized SiC particles, Compos. Sci. Technol. 68 (2008) 3033-3040. [4] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. Soc. 19 (1999) 2405-2414. F. Monteverde, L. Silvestroni, Combined effects of WC and SiC on densification and thermo-mechanical stability of ZrB2 ceramics, Mater. Des. 109 (2016) 396-407. F. Monteverde, S. Guicciardi, A. Bellosi, Advances in microstructure and mechanical properties of zirconium diboride based ceramics, Mater. Sci. Eng. A 346 (2003) 310-319. [7] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [8] W. Li, Y. Zhang, X. Zhang, C. Hong, W. Han, Thermal shock behavior of ZrB2-SiC ultra-high temperature ceramics with addition of zirconia, J. Alloy Compd. 478 (2009) 386-391. P. Hu, W. Guolin, Z. Wang, Oxidation mechanism and resistance of ZrB2-SiC composites, Corros. Sci. 51 (2009) 2724-2732. S.-Q. Guo, Densification of ZrB2-based composites and their mechanical and physical properties: a review, J. Eur. Ceram. Soc. 29 (2009) 995-1011. [11] D. Gao, Y. Zhang, C. Xu, Y. Song, X. Shi, Oxidation kinetics of hot-pressed ZrB2-SiC ceramic matrix composites, Ceram. Int. 39 (2013) 3113-3119.  [10]  [9]  [6]  5318  \\x0c']"
},{
  "_id": 229,
  "PDF": "Qualitative analysis of hafnium diboride based ultra high temperature ceramics under oxyacetylene torch testing at temperatures above 2100 degrees C.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 34 (2014) 1045-1051  Qualitative analysis of hafnium diboride based ultra high temperature ceramics under oxyacetylene torch testing at temperatures above 2100     C  Carmen Carney a,b,∗  , Anish Paul c , Saranya Venugopal c , Triplicane Parthasarathy a,b , Jon Binner c , Allan Katz a , Peter Brown d  a Air Force Research Laboratory, Wright-Patterson Air Force Base, OH 45433, United States b UES Inc., Dayton, OH 45432, United States c Department of Materials, Loughborough University, Loughborough LE11 3TU, UK d Defence Science and Technology Laboratory, Porton Down, Salisbury, Wiltshire SP4 0JQ, UK  Received 6 August 2013; received in revised form 30 October 2013; accepted 11 November 2013  Available online 5 December 2013  Abstract     Oxidation tests were carried out on HfB2 -SiC, HfB2 -HfC, HfB2 -WC-SiC, and HfB2 -WSi2 ceramics using an oxyacetylene torch. The samples were oxidized between 2100 and 2300 C. From cross-sectional images, scale non-adherence was noted as a limiting factor in oxidation resistance. The sample with the best scale adherence was HfB2 -WSi2 . Factors involving scale non-adherence such as vapor pressure, coefficient of thermal expansion mismatch and phase transformations were considered. In comparing the scale adherence of the samples it was hypothesized that vapor pressure buildup is the principal contributing factor in the scale adherence differences observed among the tested samples. However, the coefficient of thermal expansion mismatch and HfO2 phase transformation cannot be neglected as contributing factors to scale non-adherence in all samples. © 2013 Elsevier Ltd. All rights reserved.  Keywords: Hafnium diboride; Silicon carbide; Tungsten silicide; Oxidation; Ultra high temperature ceramics  1.   Introduction     Transition metal borides and carbides with melting  temperatures exceeding 2700 C are commonly  referred  to as ultra high  temperature ceramics (UHTCs) and have been studied as primary candidates  for extreme environment  thermal protection  systems  such as  those  found at  the  sharp  leading edges of hypersonic vehicles.1,2 Most  commonly  explored  are  the ZrB2 -SiC and HfB2-SiC (MeB2 -SiC) systems with and without various additives. Most oxidation resistance  testing of UHTCs has  involved either resistive-element furnace heating or arc  jet heating. Over  the past decade  the cost and  limited availability of arc  jet  testing and  the  temperature and heating  rate  limitations of furnace heating have  led many  laboratories  to develop new testing methods in order to probe higher temperatures. The first widely reported test, direct electrical resistance, developed  ∗  Corresponding author. Tel.: +1 019372556910.  E-mail addresses: carmen.carney.ctr@us.af.mil, cdoud137@yahoo.com,  ccarney@ues.com (C. Carney).  0955-2219/$ - see front matter © 2013 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2013.11.018     primarily under Halloran,3,4 provided  insight  into  the volatile nature of  the oxidation products of ZrB2-SiC at  temperatures up  to 2100 C. Observations of mixing between  the ZrO2 and SiO2 led  to a better understanding of  the dynamic characteristics occurring at  testing  temperature. Using  this  technique, oxidation resistance and mechanical strength retention comparisons between samples with different chemistries can easily and rapidly be examined.5,6 The main drawback of this method, the internal heating, has  limited  its widespread use. Laser  testing of UHTC materials has been utilized by  researchers  to  reach temperatures from benign  to beyond  the melting points of  the UHTC materials.7,8 Laser heating  technology  is versatile and more economical than arc jet testing, yet is also not widely available. As such, several laboratories have developed oxyacetylene torch  testing as a  rapid screening  tool  for UHTC materials at temperatures up  to 3400 C.9-13 The oxyacetylene  torch  test  is naturally ablative owing to the high velocity gas flow associated with  the flame. The oxidation characteristics of  the oxyacetylene  torch can be controlled by  the acetylene  to oxygen ratio, while at a given ratio  temperature  is controlled by  the distance to the flame. The torch has been used as a standard test method           \\x0c', '1046   Table 1  C. Carney et al. / Journal of the European Ceramic Society 34 (2014) 1045-1051  Sample composition and sintering parameters with sample ID.  Sample ID  HS   HSW   HH   HW   Composition   HfB2 -20vol.%SiC  HfB2 -20vol.%SiC-4vol.%WC HfB2 -4vol.%HfC HfB2 -4vol.%WSi2  Sintering temp., hold time, pressure           2100  2100  2100  1940  C, 25 min, 32 MPa  C, 15 min, 32 MPa  C, 50 min, 32 MPa  C, 15 min, 32 MPa        for oxyacetylene ablation testing of thermal insulation materials (ASTM E285-08). As UHTCs have been  tested using  these various  techniques with parameters  that  include higher  temperatures  (>1800 C) and high velocity flow,14-19 it has become apparent that there is a point at which the protective SiO2 -MeO2 scale that is formed on MeB2-SiC begins  to fail. Above a critical  temperature,  the viscosity of glassy SiO2 will be  too  low  to  remain  integral  to the scale. Under  these conditions  the SiO2 will flow out of  the pores of  the MeO2 and from  the sample surface resulting  in a less protective porous outer scale.19 Oxide scale non-adherence at temperatures in excess of 2000 C has also been observed during oxidation testing of MeB2-SiC materials.9,20,21 The loss of scale adherence could be a result of many factors including stress induced by the difference in thermal expansion coefficient (CTE) of  the MeB2 and MeO2 -based  layers upon heating/cooling, phase  transformation of  the MeO2 and  its associated volume increase  upon  cooling  or  fracture  caused  by  the  escape  of gaseous by-products of oxidation such as CO, SiO and B2O3 . The present paper  focuses on evaluating  the  relative performance  of  four  different  compositions  among  the HfB2 containing UHTCs,  at  temperatures  above 2100 C  for  long duration, using an oxyacetylene  torch. A baseline  sample of HfB2-SiC was  tested and compared  to W-containing samples. W has been  found  to be a beneficial additive  for both ZrB2 and HfB2 -based UHTCs when  samples are  tested under  furnace heating up to 2000 C,20,22,23 and as such two chemistries including W, HfB2 -WC-SiC and HfB2-WSi2 , were  tested. A qualitative analysis of the differences in behavior is presented.        2. Materials and methods  Commercially  available  HfB2 (Materion,  99.9%, −325 mesh), WC (Materion, 99.5%,  −325 mesh), WSi2 (Mate−325 mesh), Hf (Materion, 99.8%,  −325 mesh), rion, 99.5%,  −325 mesh) and  C (Materion, 99.5%,  ␤-SiC (Materion, 99.9%, 1  \\u242em) were used  to prepare  four different  sample  compositions: HfB2-20vol.%SiC [HS], HfB2-20vol.%SiC-4vol.%WC [HSW], HfB2-4vol.%HfC [HH] and HfB2-4vol.%WSi2 [HW]. The powder mixtures were ball milled  in  isopropanol for 24 h with SiC grinding media, dried at room temperature, and subsequently dry milled for 12 h. Typical weight loss of the SiC grinding media after milling was 0.2 mg (0.2 wt.% of the total batch). The powders were sieved through an 80-mesh (177  \\u242em) screen. Sample composition and  sintering conditions are  summarized  in Table 1. Milled powders were  loaded  into a 20 mm diameter graphite die to produce a 13 mm thick cylinder. A layer of BN and graphite foil separated the powder from the die with                 the powder  in contact with  the graphite foil. The powder-filled dies were cold pressed at approximately 50 MPa. The powders were sintered using field assisted sintering (FAS: FCT Systeme GmbH, Model HPD 25-1, Rauenstein Germany) at 2100 C (HSW, HH, HW) or 1940 C  (HW)  for 25 min  (HS), 15 min (HSW and HW), or 50 min (HH) under a 32 MPa load. The controlled heating and cooling rates were 50 C/min. The load was applied during heating  to 1600 C and  released on cooling  to 1000 C. The graphite  foil was  removed  from  the sample by manual grinding, and  the  faces of  the disks were polished  to 45  \\u242em for exposure to the torch. Samples were oxidized using  the oxyacetylene  torch apparatus developed at Loughborough University.13 Samples were held  in place using a carbon-carbon  foam  insert  in a watercooled graphite holder. The surface of  the sample was 25 mm from  the exit point of  the  torch. During  the  test,  the  temperature of the exposed face was recorded using a 2 color pyrometer (Marathon MR1SCSF, Raytek GmbH, Berlin, Germany) while the back  face  temperature was measured by a J-type  thermocouple inserted through a hole in the graphite holder to contact the back of  the sample. Fig. 1 shows a photograph of  the  test along with  the  front  face heating profile  for  the samples. The photograph (Fig. 1a) was taken through welding glasses to limit the brightness. An oxygen  rich flame was chosen  for  testing the UHTC samples. The acetylene  −1 and 1.1 m3 h to oxygen ratio was 1:1.35 −1 respectively. The with flow rates of 0.8 m3 h heat flux was measured at 25 mm using  the same acetylene  to oxygen flow rates and ratio with a water-cooled gardon gauge (TG1000-54, Vattel Corp., Christiansburg, VA). Oxidized samples were analyzed by X-ray diffraction (XRD: D8 Bruker AXS limited, Coventry, UK), and  then mounted  in epoxy and cut in half. The cross section of  the oxidized  face was  then pol\\u242em. The microstructures were characterized using ished  to 1  scanning electron microcopy (SEM: Quanta, FEI, Hillsborough, OR) along with energy dispersive spectroscopy (EDS: Pegasus 4000, EDAX, Mahwah, NJ) for elemental analysis. Wavelength dispersive spectroscopy  (WDS: SX100, Cameca, France) was used for the chemical analysis of W and Si. Transmission electron microscopy (TEM:Phillips CM200 FEI, Hillsborough, OR) samples were prepared by  focused  ion beam milling  (FIB: DB235, FEI, Hillsborough, OR).  3. Results  3.1. Oxidation test parameters  The oxyacetylene  monality  in use, but   torch  test  is one  that  is gaining  comis not completely characterized. Sample  \\x0c', 'C. Carney et al. / Journal of the European Ceramic Society 34 (2014) 1045-1051   1047  The pyrometer data recorded during sample oxidation is plotted in Fig. 1b. Tests were targeted for a 8 min duration and this was achieved for all but the HS sample, which was liberated from a degraded holder after 3.6 min. The samples were positioned 25 mm from the torch nozzle after it had been lit. Transient heat up time was about 90 s. The samples reached peak temperatures between 2100 and 2300 C. The average temperature value noted in Fig. 1b  is  the average of  the  last 60 s of exposure. The back face  temperature  reached approximately 1200 C  for  the HW and HS samples and held steady during  the hold  temperature. The challenge in measuring the back face temperature with this method was shown by the instability of the back face temperature for the HS and HH samples which fluctuated between 800 and 1300 C during the hold temperature.           3.2. Analysis of the oxide scales     Optical photographs of  the  sample  surfaces are  shown  in Fig. 2. The hottest part of  the flame was  slightly off-center creating  a  cooler  crescent-shaped  region  (evidenced by  less extensive oxidation) on the sample surface. A temperature variation  is present on  the surface of  the sample due  to  the flame being smaller  than  the diameter of  the sample. A difference of 300-400 C was observed from  the center of  the sample  to  the outer diameter during prior testing of 15 mm diameter samples of similar composition  that were placed 20 mm  from a flame with  the same flow and composition. Evidence of oxide nonadherence was observed  in each sample. The oxide scale was white  in all samples except  the HSW and HW samples, which had areas colored  light green  that could be evidence of WO3 . XRD analysis (Fig. 3) of the oxidized sample surfaces showed that  in all cases monoclinic HfO2 was  the primary crystalline constituent. In  the HS and HSW samples, peaks of HfB2 were observed; probably as a result of the cracks in the oxide exposing the underlying bulk material. The HSW sample contained a peak near 40 , not seen in any of the other samples. Considering the sample chemistry,  the peak could be attributable  to various W-containing species such as W or WxSiy . SEM analyses of  the sample cross sections across  the center of the hot zones clearly show the non-adherent nature of the oxide scales  (Fig. 4). The  images  in Fig. 4 are  from  the center of  the sample. Samples HS  (Fig. 4a), HSW  (Fig. 4b), and HH  (Fig. 4d) exhibit multiple oxide  layers, with each subsequent oxide scale  forming on surfaces below which  the scale had previously disadhered. Adherence was  limited  to  the sample perimeter, where  the scale was  thinner and  the sample has     Fig. 1.   (a) Photograph of   the   torch heating process. The sample   is moved   in  manually from a position out of the flame to a position 25 mm from the flame. A  is the water-cooled graphite holder, B is the C/C insert to hold the sample in place  and C   is   the sample. (b) Plot of   the measured pyrometer   temperature readings  during oxyacetylene torch testing of the HfB2 -based samples. Temperature was measured at  the center of  the front face of  the sample and  is plotted from  the  beginning of the test until the torch was extinguished. The average temperature  was calculated from the last 60 s of exposure.  temperature  is dictated by sample composition (its heat capacity  and  thermal  conductivity),  the  sample distance  from  the flame, and  the oxygen  to acetylene ratio. In addition,  the oxyacetylene flame environment is complicated by the presence of carbon  species  that depend on  the flame chemistry.24,25 The heat flux can be measured by calorimetry. When the flow rates, gas  ratio and distance  to  the flame are  set,  the heat flux as measured by  the calorimetry  is fixed. The heat flux was mea−2 at 25 mm from the torch exit using an sured to be 880 W cm acetylene to oxygen ratio of 1:1.35 with flow rates of 0.8 m3 h −1 respectively. and 1.1 m3 h  −1  Fig. 2. Photographs taken after oxyacetylene torch testing of the HfB2 -based UHTCs.  \\x0c', '1048   C. Carney et al. / Journal of the European Ceramic Society 34 (2014) 1045-1051     absorption coefficient of W and the surrounding oxide, a precise phase  identification was not achieved. A W-phase with similar morphology has also been observed in furnace and direct electrical  resistance heating of W-containing HfB2 -based samples C and above.20,23 heated to 2000 The oxide scale in the HW sample (Figs. 4c and 5c) was distinctly different. Although some separation between  the scale and bulk sample was observed, areas of adhered scale were also found near  the center of  the sample. Additionally,  the regions of non-adherent oxide were a single layer that broke away from the underlying bulk. A comparison of the W-containing phases in HSW and HW are indicated in Fig. 6. In the HW sample, the W-phases were present in the denser outer oxide scale (Fig. 6(b)) and in the more porous inner scale (Fig. 6(c)), while W-phases were only found in the interior porous HfO2 of the oxidized HSW sample. A Cu  impurity  in  the WSi2 (0.001 vol.%, as reported) manifests as Cu found with W at  the bulk-oxide  interface and in some  instances with  the W-phase  in  the oxide scale. After annealing at 200 C,  thin films of 45  to 72 at.% Cu  in W can be formed, but at room  temperature  the systems show negligible mutual solubility.27 EDS estimates place  the Cu content at 7-9 at.% if a Cu-W alloy is assumed. However, the low symmetry electron diffraction patterns cannot be matched  to a W-Cu solid solution. As shown by TEM  (Fig. 6(d)) a W-containing grain boundary phase exists between  the  large W-containing phases and HfO2 in  the oxide scale. The volume of  this phase precludes its identification.     4. Discussion     ×  The common aspect of  the oxidation behavior among HS, HSW and HH  is  the formation of multiple  layers, presumably from repeated separation of scale from the bulk. There are at least three possible reasons why  the oxide scale may separate from the bulk sample during testing: (1) vapor pressure buildup at the scale-bulk interface; (2) CTE mismatch between the oxide scale and  the bulk; and (3)  the phase  transformation between monoclinic and tetragonal HfO2 and its associated volume change. When a vapor species or combination of vapor species at the bulk/oxide  interface  exceeds 1 atm  (1   105 Pa) pressure the oxide scale may be disrupted. Thermodynamic calculations made by Opeka et al.28 and Fahrenholtz29,30 on  the ZrB2-SiC system at 2227 C reveal that the vapor pressure of Zr (and Hf) oxides are at  least 4 orders of magnitude  lower  than  the 1 atm limit. It has been shown that B2O3 does exceed the 1 atm criteria at 1950 C,28 but  its vapour pressure  is sufficiently high above 1200 C  to cause evaporation from  the surface  long before  the internal pressures exceed 1 atm. The SiO partial pressures can exceed 1 atm at  the  interface at  temperatures above 1865 C, while the vapour pressure of CO is much higher than this. The dissimilarity in the appearance of the HW oxide scale suggests a different set of mechanisms are at work. Unlike  in  the HS, HSW and HH samples, there are no C-containing phases in the HW sample, suggesting that CO evolution may play a role in the lack of oxide adherence during testing. This mechanism would be active during heating of the sample.           Fig. 3. XRD plot of the four HfB2 -based UHTC monoliths after oxyacetylene torch  testing. All peaks are HfO2 except  the peaks  labeled *  (HfB2 ) and  (possibly a W-phase).    Fig. 4. SEM micrographs of an overview of the oxide scale formed on (a) HS,  (b) HSW,   (c) HW, and   (d) HH after oxyacetylene   torch   testing. Spallation   is  evident   in all samples with multiple   layers forming on   the HS, HSW, and HH  samples. The dark continuous phase is mounting epoxy while the light regions  correspond   to   (O)   the oxide scale,   (B)   the bulk, unoxidized sample, and   (D)  SiC-depleted layer.  been shown  in previous experiments  to be cooler by a  thermal imaging camera. Fig. 4 shows an overview of the oxide scale of each sample. Magnified  images of  the oxide scales are shown in Fig. 5. The outer oxide layers were HfO2 in the HS and HH samples, while HfO2 and a W-phase were observed in the HSW sample. Oxide scale  layers between  the outermost oxide  layer and  the bulk of  the HS and HSW samples retained SiO2 glass while SiO2 could also be found at the interface between the nonadherent layers and the bulk sample. SiO2 was likely removed by the flowing gases as velocities up to 70 m/s have been reported at distances 15-30 mm from an oxyacetylene cutting torch.26 EDS of  the W-phase found with HfO2 in  the HSW sample showed primarily W, but because of the small size of the phase, the high  \\x0c', 'C. Carney et al. / Journal of the European Ceramic Society 34 (2014) 1045-1051   1049  Fig. 5. SEM micrographs of a magnified view of   the oxide scale   formed on   (a) HS,   (b) HSW,   (c) HW, and   (d) HH after oxyacetylene   torch   testing. Labeled  regions are (a) A = primarily HfO2 and B = HfO2 /SiO2 ; (b) C = primarily HfO2 and D = HfO2 /W-phase; (c) E = dense HfO2 /W-phase, F = porous HfO2 /W-phase and G = HfB2 /W-phase/Si-phase; and (d) H = HfO2 .           C.2,6,10,21 But   In addition, prior studies of SiC oxidation report bubble formation within the SiO2 scale at 1700 C, attributed to CO partial pressures up  to 7.5 atm.31 Luthra32 argues for a SiC oxidation mechanism  that  involves both diffusion  limited and  interface limited reactions where CO gas bubbles can form if the permeabilities or diffusion  rates of CO are substantially  lower  than that of oxygen or the SiC or bulk-oxide interface is C-rich. Such C-rich deposits have been observed at  the bulk-oxide  interface in oxidized MeB2 -SiC systems.21,33 Bubbling of the SiO2 scale has also been reported in MeB2-SiC systems between 1500 and 2200 2300 it  is possible  that for HfB2 -based systems heated  to  C, sintering of  the outer HfO2 could hinder vapor escape and lead to spallation as a result of the buildup of vapor. In fact, sintering of the outer HfO2 scale is suggested in Fig. 5 and has been reported in ZrB2 -based oxidation studies at temperatures in excess of 1800 Previously HfB2-SiC and HfB2 -SiC-WC have been  tested by  direct  electrical  resistance.20 The  samples  heated  to 2000 C  by  direct  electrical  heating with  a  heating  rate of approximately 2000 C/min  showed oxide  scale  spallation between  the bulk unoxidized material and  the oxide scales  in the HfB2-SiC20 and HfB2-SiC-WC  both  (unpublished) samples. The location of the spallation suggests that the coefficient of  thermal expansion difference between  the  scale and bulk may play a  role  in  scale adherence. Another mechanism by which stress would be  introduced at  the  interface between  the oxide scale and bulk sample  is  through  the phase  transformation and reverse  transformation from monoclinic  to  tetragonal  C.10,34              HfO2 during heating and cooling. No tetragonal phase stabilization was found in either the HSW or HW sample, as expected.35 As for CTE modification, hafnium and zirconium tungstates are known  to have a negative CTE, but are only stable  in a narrow temperature regime, 1276-1105 C for HfW2O8 .36,37 The latter can be quenched to room temperature with fast cooling rates, but the byproducts are HfO2 and WO3 .36 if decomposition occurs  It is possible with the rapid cooling experienced by the samples some HfW2O8 remains. Although  the phase was not observed by XRD,  the concentration could be below  the detection  limit or sufficiently deep within  the oxide  to avoid detection. EDS scans of  the W-containing phases  from  the dense outer oxide scale  in  the HW oxide scale show  the presence of W and Cu. A grain boundary phase  is  found between  the  large W-phase and HfO2 grains. This phase has not been previously reported or observed  in oxidized W-containing HfB2 samples. Because of  the small volume of  the phase a precise composition could not be determined. A WO3-HfO2 liquid phase can form at temperatures above 1280 C,38 suggesting  that such a phase could form during testing. Determination of the mechanisms leading to non-adherence may allow engineering of UHTC compositions that form more resilient oxide scales. Further  information  regarding chemical composition of the oxide scale and whether the oxide scale spalls during  the hold at  testing  temperatures (most  likely  to be as a result of vapor pressure effects) or upon cooling (corresponding to CTE mismatch or phase transformations) could improve this understanding.     \\x0c', '1050   C. Carney et al. / Journal of the European Ceramic Society 34 (2014) 1045-1051  Fig. 6. SEM micrograph of the W-containing phases (lighter) found with HfO2 after oxyacetylene torch testing of (a) HSW and (b and c) HW. The phases in (b) are found at the porous inner oxide, while the phases in (c) are found in the dense outer oxide of the HW sample. The W-phase in (a) and (b) are similar in morphology  and chemistry (see   inset EDS   in (b)). The W-phases found   in   the dense   layer of   the HW sample are   larger and   in some cases contain Cu. The C   in   the EDS   is an  artifact of   the C-coating applied for SEM analysis. WDS (see   inset WDS   in (c)) confirms W and Cu without Si. (d) TEM micrograph of a W-containing phase   in  HW from the dense outer oxide scale. Phase 1 was shown by TEM-EDS to contain W and Cu, while phase 2 is a grain boundary phase potentially composed of Hf,  W, and O.  5. Conclusion  Oxyacetylene torch testing is an aggressive test for assessing the oxidation and ablative resistance of ultra high  temperature ceramics. All samples formed outer HfO2 -based oxide scales. In the HS, HSW, and HH samples, the oxide scale consisted of multiple  layers  that were non-adherent  to  the underlying bulk material. From the multiple layers and the presence of new layers forming on the non-oxidized bulk, it is possible that as each layer broke away from the surface oxygen was able to penetrate and form new oxide layers. The HW sample was distinct in that its oxide scale consisted of a single, partially adherent layer. The differences in the samples suggest that the evolution of CO during oxidation may adversely impact scale adherence. However, the CTE mismatch and HfO2 phase  transformation cannot be ruled out as  important factors  in establishing adherence of  the oxide scales.  Acknowledgements  The oxyacetylene torch testing was performed under a Project Arrangement between  the United States of America Department of Defense and the United Kingdom of Great Britain and  Northern  Ireland Ministry of Defence under  the guidance of Dr. Allan Katz (AFRL/RX) and Professor Peter Brown (Dstl).  References  1. Wuchina   E, Opeka M,   Causey   S,   Buesking K,   Spain   J,   Cull A,  et   al. Designing   for ultrahigh-temperature   and  thermal  properties  2004;39(19):395939-49.  of HfB2 , HfCx , HfNx  and   applications:The mechanical ␣Hf.   J Mater   Sci     2. Opeka MM, Talmy IG, Zaykoski JA. Oxidation-based materials selection for  2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci 2004;39(19):5887-904.  3. Karlsdottir SN, Halloran JW. Rapid oxidation characterization of ultra-high temperature ceramics. J Am Ceram Soc 2007;90(10):3233-8.  4. Karlsdottir SN, Halloran JW, Henderson CE. 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},{
  "_id": 230,
  "PDF": "Quantitative Electron Microprobe Characterizations of Oxidized ZrB2 Containing MoSi2 Additives.pdf",
  "Text": "['Oxid Met (2009) 72:335-345  DOI 10.1007/s11085-009-9165-y  C O M M U N I C A T I O N  Quantitative Electron Microprobe Characterizations  of Oxidized ZrB2 Containing MoSi2 Additives  Shu-Qi Guo Æ Takashi Mizuguchi  Æ Takeshi Aoyagi  Æ  Takashi Kimura Æ Yutaka Kagawa  Received: 27 March 2009 / Revised: 21 June 2009 / Published online: 16 July 2009 Ó Springer Science+Business Media, LLC 2009  Abstract  This study reports the oxidation behavior of hot-pressed ZrB2 ceramics, with 10, 20 and 40 vol.% MoSi2 additives, exposed to dry air at 1500 °C for up to  10 h. The effect of the amount added MoSi2 on oxidation resistance was assessed.  Quantitative electron microprobe characterizations of the oxidized ZrB2 with MoSi2  additives were carried out with electron probe micro analysis. Parabolic oxidation  behavior was observed for  the three compositions. The oxidation resistance was  significantly improved with MoSi2 additives. Reaction-product phase compositions  and phase distribution were  thoroughly characterized from the oxidized surface  through to the unreactive bulk material.  It was  found that  the oxidation products  consisted of nonstoichiometric  amorphous SiO2,  stoichiometric  crystalline ZrO2,  and MoB. The  amounts  of  these  phases  in  the  oxidized  reactive  region were  qualitatively determined.  Keywords  Zirconium diboride \\x01 Molybdenum disilicide \\x01 Oxidation \\x01 Quantitative characterization \\x01 EPMA  S.-Q. Guo (&)  \\x01 T. Mizuguchi  \\x01 Y. Kagawa  Hybrid Materials Center, National Institute for Materials Science, 1-2-1 Sengen,  Tsukuba 305-0047, Ibaraki, Japan  e-mail: GUO.Shuqi@nims.go.jp  T. Aoyagi  \\x01 T. Kimura  Advanced Nano Characterization Center, National Institute for Materials Science,  1-2-1 Sengen, Tsukuba 305-0047, Ibaraki, Japan  Y. Kagawa  Research Center for Advanced Science and Technology, The University of Tokyo,  4-6-1 Komaba, Meguro-ku, Tokyo 153-8505, Japan  123  \\x0c', 'Introduction  Zirconium diboride (ZrB2)-based ceramic materials have an extremely high melting ([3000 °C), high thermal  point  and electrical  conductivities,  chemical  inertness  against molten metals,  and good thermal  shock resistance  [1, 2]. These unique  mechanical  and physical properties have never been achieved by other  ceramic  materials. Currently, ZrB2-based ceramics are considered potential candidates for a  variety  of  high-temperature  structural  applications,  including  furnace  elements,  plasma-arc electrodes, rocket engines and thermal protection structures for leadingedge parts of hypersonic re-entry space vehicles operating at over 1800 °C [3-5].  For  these  applications,  a major  problem with  ZrB2-based  ceramics  is  high  temperature  oxidation  [6-8].  For  example,  pure  ZrB2  is  easily  oxidized  in  atmospheric air to form ZrO2 and B2O3. Among these oxidation products, B2O3 has a high vapor pressure and is vaporized above 1300 °C. Thus,  the oxidation  resistance of ZrB2 must be improved for use in oxidizing environments.  To improve oxidation resistance, Si-containing additives, such as SiC and MoSi2,  are made to ZrB2 [8-13] in order to form a protective borosilicate glass temperatures above 1200 °C that enhances  layer at  the oxidation resistance of ZrB2. The  oxidation behavior of ZrB2 with the SiC and MoSi2 additives has been examined  previously [8-11]. These studies demonstrated that  the oxidized reactions of ZrB2  ceramics depended on the composition, exposure temperature, exposure time at a  particular temperature, and oxygen content  in the oxidizing atmosphere. Hence,  the  oxidation  scale  is  a  complex  one,  consisting  of  varying  compositions  and  morphologies from the top surface to the unreactive bulk material. For the ZrB2- the main oxidation products are ZrO2, B2O3, and SiO2 below 1300 °C, above 1300 °C,  SiC system,  and ZrO2  and SiO2  as B2O3  liquid gets  completely vaporized.  Generally,  the resulting oxidized reactive  region consists of  three layers:  (i)  the  outermost glassy layer  (SiO2),  (ii)  the oxide subsurface layer  (ZrO2 ? SiO2), and  (iii)  the  SiC-depleted ZrB2  layer  (ZrB2 ? SiC ? ZrO2).  For  the ZrB2-MoSi2  system, on the other hand, MoSi2 is oxidized in atmospheric air, instead of SiC. Sciti  et al. [12, 13] examined the oxidation behavior of pressurelessly sintered 20 vol.% MoSi2-containing ZrB2 between 700 °C and 1400 °C in dry air. They showed that  the resistance to oxidation significantly improves by the addition of MoSi2, because  of the SiO2 scale that forms on the surface. The oxidation products consist of SiO2,  ZrO2, ZrSiO4, MoO2, and MoB, and the oxidized reactive region consists of a SiO2 rich glass layer; a subsurface oxide layer; and a MoB, ZrO2, and SiO2-containing  mixture layer, depending on the exposure temperature. Additionally, although the  oxidation products and oxidized layers of these composites have been characterized  by X-ray  diffraction  (XRD),  electron  scanning microscopy  (SEM),  and  energy  dispersive  spectroscopy  (EDX),  detailed  information  about  the  composition,  distribution, and amount of  reactive phases in the oxidized reactive region is not  well known because of  the limitations of  these methods. Therefore, quantitative  characterization  of  oxidized  reactive  regions  is  required  for  understanding  the  oxidation behavior/mechanism. Additionally,  the effect of the amount of MoSi2 on  oxidation resistance needs to be explored. In this study, hot-pressed ZrB2 ceramics with 10, 20, and 40 vol.% MoSi2 additives were exposed in dry air at 1500 °C for  336  Oxid Met (2009) 72:335-345  123  \\x0c', 'up  to  10 h. Quantitative  characterization  of  oxidized ZrB2 with MoSi2 were  conducted using electron probe micro analysis  (EPMA)  to identify the oxidation  products,  reactive phase composition, amount, and distribution. Also,  the effect of  the amount of MoSi2 on oxidation resistance was examined.  Experimental  The material used in this  study was prepared by hot-pressing MoSi2-containing  ZrB2  ceramics.  In order  to examine  the  effect of  the  amount of MoSi2 on the  oxidation resistance of ZrB2,  three compositions of 10, 20, and 40 vol.% MoSi2 modified ZrB2  ceramics were  sintered. The detailed sintering process has been  reported elsewhere [14]. Hereafter,  the three compositions are denoted as ZM10,  ZM20, and ZM40. Test  specimens with average dimensions of 5 9 2.5 9 2 mm  were cut from the hot-pressed ZrB2 ceramic plates using a conventional mechanical  cutting  procedure. After  all  the  surfaces were  polished,  these  specimens were  ultrasonically cleaned in acetone and then kept in an oven at a constant temperature of 100 °C prior to oxidation. Oxidation tests were performed using an electronic furnace (BFT-150-P, Nikkato Co., Ltd., Tokyo, Japan) at 1500 °C over a period of 10 h in dry air. The heating and cooling rates were 20 and 10 °C/min, respectively.  Before and after oxidation, the specimens were weighed using an analytical balance  (AB265-S, Mettler Toledo Co., Ltd., Switzerland) with an accuracy of 0.1 mg. The  surface morphologies of specimens were observed by SEM, combined with EDX,  with an accelerating voltage of 15 kV.  In order to examine the oxidation evolution, the oxidized specimens were cut into  two parts  from the  central  section of  the  specimens,  and one of  the parts was  mounted in epoxy and then carefully polished. Observation of the cross sections of  the oxidized samples’ microstructures were conducted using a prototype wavelength  dispersive-electron  probe microanalyzer  (WDS-EPMA)  with  an  accelerating  voltage  of  20 kV, which was  developed  by Kimura  et  al.  [15]  based  on  a  commercially available FE-EPMA (JEOL; JXA8900R). In addition, map analysis of  O, Si, Mo, Zr, and B elements within the cross section plane was conducted at 0.2 lm/step in the pixels of 200 9 1024 using an X-ray mapping technique with EPMA, with an accelerating voltage of 10 kV and probe current of 5.0 9 10-8 A.  The X-ray image data obtained is projected on a two-dimensional space of X-ray  intensity to form a histogram, which is generally called a scatter diagram. Detailed  calculations of the scatter diagram have been reported elsewhere [16]. The analysis  of  the scatter diagram,  i.e.,  the scatter diagram method, was used to identify the  compounds  composition  as well  as  determine  their  distribution, with  the  pure  stoichiometric SiO2, ZrO2, MoB, ZrB2 and MoSi2 powders as reference materials.  Furthermore,  the EPMA micrographs were used for quantitative analysis by means  of an image processing system (LUZEX III, Nireko Co., Ltd., Tokyo, Japan). The  volume  fraction  of  each  reactive  phase  in  the  oxidation  reactive  region was  determined.  Oxid Met (2009) 72:335-345  337  123  \\x0c', 'Results and Discussion  In Fig. 1, the plots of weight gain as a function of exposure time for the hot-pressed MoSi2-containing ZrB2 oxidized at 1500 °C are presented. The three compositions  of ZrB2  ceramics  show similar  oxidation  behaviors:  the  specific weight  gain  increases rapidly at  the initial stage of exposure (within 2 h) and then the weight  gain increases gradually with exposure time. The three compositions have specific weight gains of the order of 6.5-23.1 mg/cm2 after 10 h of oxidation exposure at 1500 °C, depending on the amount of MoSi2. The specific weight gains of the three  compositions decrease with the amount 10 [ 20 [ 40 vol.%. This suggests that  of MoSi2  added,  in  the  order  increasing MoSi2 is effective in improving  the oxidation resistance of ZrB2.  Figure 2 presents plots of the square of the weight gain, W2, as a function of three compositions after oxidation at 1500 °C. The straight  oxidation time,  t,  for  lines closely represent parabolic oxidation kinetics  for each curve. This  suggests  that  the diffusion,  such as outward diffusion of  additive  cations  from the bulk  material  to the oxidized surface and the inward diffusion of oxygen through the  oxide  layer,  is  the  rate-controlling mechanism for  oxidation.  The  parabolic  oxidation rate constants of three compositions, obtained from the slopes of straight lines, are 54.4 mg2 cm-4 h-1 for ZM10, 13.2 mg2 cm-4 h-1 for ZM20, and 4.4 mg2 cm-4 h-1 for ZM40. However, additional data corresponding to longer  the  oxidation time may be needed to determine whether  the oxidation behavior of  the  studied material  is parabolic or not.  Figure 3 shows XRD patterns of three compositions before and after 10 h of oxidation exposure at 1500 °C. Note that  for pristine ceramic materials, only the  XRD pattern of ZM20 is shown because the same phases are detected for the ZM10  and ZM40 compositions, except  for  the different  intensities in their peaks. Before  oxidation  exposure,  only ZrB2  and MoSi2  are  detected  for  each  composition  (Fig. 3a). After oxidation (Fig. 3b-d), a new primary oxidized phase of ZrO2  is  detected for  the three compositions. The intensity of  the peaks of  the ZrO2 phase  decrease with the amount of MoSi2 added, showing improved oxidation resistance.  This agrees with the decrease of weight gain observed with the increasing amount of  ZM10  1500°C  (  m  g  /  c  m  2  )  t t  G  a  i  n  ,  W  ZM20  W  e  i  h g  ZM40  Time at Temperature t (h)  25 25  30  20  10 10  15  5  0  0  2  4  6  8  10  12  Fig. 1  Plots of weight gain as a  function of exposure time for  hot-pressed ZrB2 with MoSi2 oxidized at 1500 °C  338  Oxid Met (2009) 72:335-345  123        \\x0c', 'MoSi2 additive (Figs. 1 and 2). For comparison,  the intensity of  the ZrB2 peaks is  higher  for ZM20 and ZM40 than for ZM10,  indicating that  those signals are from  the bulk ceramic beneath the oxide layer. Additionally,  the MoB phase is detected  for both ZM20 and ZM40, but the peaks of the MoSi2 phase are not detected in any  instance. These facts show that MoSi2 is oxidized during exposure. In particular, the  absence  of MoSi2  peaks  indicates  that MoSi2  beneath  the  oxide  layer  is  so  thoroughly oxidized during exposure that it can-not be detected by XRD. Generally,  oxidation of MoSi2 results in the formation of SiO2 on the surfaces of specimens.  However,  the SiO2  peak  is  not  detected  in  any  instance  after  oxidation. This  suggests that although the peak of SiO2 is absent in the XRD profile, the SiO2 phase  formed  on  the  surface  of  specimens  during  exposure  should  be  present  in  an  amorphous form.  600 600  700  ZM10  1500°C  500  ( (  m  g  2  /  c  m  4  )  300  400  G G  a  i  n  ,  W  2  100 100  200  W  e  i  h g  t  ZM40 ZM40  ZM20  0  0  2  4  6  8  10  12  Time at Temperature, t (h)   Fig. 2  Parabolic plots of  specific weight gain as a  function of time for hot-pressed  ZrB2 with MoSi2 oxidized at 1500 °C  (d)  ZrB2  MoSi2  ZrO2  MoB  u u  .  )  ZM20  ZM40  (c)  n e e  s  i  t  y  (  a  .  ZM10  (b)  (a)  I  n  t  ZM20  80  20  30  40  50  60  70  2θ(degree)  Fig. 3 X-ray diffraction patterns of the specimen surfaces for hot-pressed ZrB2 with MoSi2; (a) before, and (b), (c), (d) after 10 h of oxidation exposure at 1500 °C  Oxid Met (2009) 72:335-345  339  123        \\x0c', 'Figure 4 shows the morphologies of the surfaces for ZrB2 with the MoSi2 additive before and after oxidation exposure of 10 h at 1500 °C. Before oxidation  (Fig. 4a),  the microstructure of ZrB2 with the MoSi2 additive consists of equiaxed  ZrB2 (brighter contrast) and MoSi2 (dark contrast) grains. After oxidation (Fig. 4b),  the  surface  is  covered with  a  glassy  layer  in which white  particles  of  \\x0c', 'Oxid Met (2009) 72:335-345  341  Fig. 4  Examples of SEM images of surface morphologies for hot-pressed ZrB2 with MoSi2; a before, b after oxidation of 10 h at 1500 °C, and c, d EDX spectra corresponding to A and B marked in b,  respectively  123  \\x0c', '342  Oxid Met (2009) 72:335-345  Fig. 5  Examples of a EPMA backscattered image of the cross section and b elemental mappings under EPMA for hot-pressed ZrB2 with MoSi2 oxidized at 1500 °C for 10 h  complex  compounds  and  consists  of  Si-O, Mo-B,  Zr-O,  and  Zr-B phases.  Furthermore,  layer  II  is divided into two sublayers:  II(a) and II(b). Sublayer  II(a)  consists of Si-O, Mo-B, and Zr-O phases,  in which Zr-O is the primary phase. In  sublayer II(b),  the Zr-B phase is also present. However,  the Mo-Si phase is absent  in the entire reactive region. This result agrees with the XRD profile in which the  MoSi2  peak  is  absent. Additionally,  in  order  to  identify  the  reactive  phase  compositions as well as  to examine a change of  the chemical compounds during  exposure,  the analysis of X-ray image data for Si, O, Zr, Mo, and B elements  is  conducted on the cross section. The scatter diagrams of the obtained characteristic  X-ray intensities  for Si-O, Zr-O, Mo-B, Mo-Si, and Zr-B phases are shown in  123  \\x0c', 'Oxid Met (2009) 72:335-345  343  Fig. 6  Examples of a X-ray image of phase mapping of the cross section under EPMA, and the scatter  diagrams of b Si vs. O, c Zr vs. O, d Mo vs. B, e Mo vs. Si, and f Zr vs. B for hot-pressed ZrB2 with MoSi2 oxidized at 1500 °C for 10 h  Fig. 6b,  c, d,  e,  and f,  in which each point  represents  the number of  the  same  intensities. Note that the solid circle in each figure represents the pure stoichiometric  SiO2, ZrO2, MoB, MoSi2, and ZrB2 phases.  It  is  found that  the Zr-O and M-B  phases produced during oxidation exposure are the stoichiometric ZrO2 and MoB  phases. For comparison,  the Si-O phase deviates from stoichiometric SiO2, with a  chemical form of Si1-xO2. The composition of the Si-O phase is determined to be  31.9 mass% Si and 51.5 mass% O. For comparison,  the stoichiometric SiO2 phase  consists of 46.74 mass% Si and 53.26 mass% O. Obviously,  the Si amount  is not  large enough to form stoichiometric SiO2 during oxidation exposure. This means  that the inward diffusion of oxygen from the oxidized surface to the bulk material is  more rapid than the outward diffusion of silicon from the bulk region to the oxidized  surface through the oxidized layer. This can be attributed to the oxidized layer of  ZrO2 (layer II(a)) being the best for fast oxygen ion conduction [18]. Thus, although  stoichiometric SiO2  is generally present  in high-temperature oxidation,  for ZrB2  with MoSi2  investigated  in  this  study,  the  Si-O phase  is  a  nonstoichiometric  amorphous  SiO2  phase. This  discrepancy  is  not  yet  fully  understood,  but  is  associated with the deficiency of Si.  Moreover, analysis of the scatter diagram of Zr-B shows that  the Zr-B phase is  stoichiometric ZrB2  (Fig. 6f). For  comparison,  the Mo-Si phase  is nonstoichio metric MoSi2 (Fig. 6e), with a chemical  form of MoSi2-x. This decrease of Si  is  attributed to the outward diffusion of Si during exposure. Thus,  it is reasonable that  the oxidation of ZrB2 with the MoSi2 additive is controlled by the inward diffusion  of O and  outward  diffusion  of  Si  and  the  oxidation  products  consist  of  the  123  \\x0c', '344  Oxid Met (2009) 72:335-345  nonstoichiometric amorphous SiO2 phase and the stoichiometric crystalline ZrO2  and MoB phases. Furthermore,  it  is found that  the volume fraction of Si-O, Zr-O,  M-B, and Zr-B phases determined within layer II from the EMPA phase mapping (Fig. 6a) is equal to *16% for Si-O, *55% for Zr-O, *18% for Mo-B, and *11% for Zr-B. To the author’s knowledge, for the first time, the reactive phases in  the reactive region have been quantitatively characterized in oxidized ZrB2 with  MoSi2 additives, with the scatter diagram method under EMPA.  Summary  In conclusion,  the oxidation resistance of ZrB2 ceramics with MoSi2 additives  is  governed by the inward diffusion of oxygen from the oxidized surface to the bulk  material and the outward diffusion of silicon from the bulk region to the oxidized  surface. The microstructure of oxidized ZrB2 with MoSi2 additive consists of an  outermost  surface dense glassy layer,  a middle oxidized reactive  layer,  and an  unreactive bulk material. The scatter diagram method with EPMA is an effective  quantitative method to characterize  the  reactive phase  compositions,  and phase  distribution  in  the  oxidized  reactive  region  from the  outermost  surface  to  the  unreactive  bulk material. The  oxidation  products  consist  of  nonstoichiometric  amorphous SiO2, stoichiometric crystalline ZrO2, and MoB. The outermost glassy  layer is nonstoichiometric amorphous SiO2;  the middle reactive layer is composed  of  complex  compounds,  nonstoichiometric  amorphous SiO2,  and  stoichiometric  crystalline ZrO2, MoB,  and ZrB2;  and  the  unreative  bulk material  consists  of  stoichiometric ZrB2 and nonstoichiometric MoSi2.  References  1. C. Mroz, American Ceramics Society Bulletin 73, 141 (1994).  2. R. Telle, L. S. Sigl, and K. Takagi,  in Handbook of Ceramic Hard Materials, ed. R. Riedel (Wiley VCH, Wiinheim, Germany, 2000), p. 803.  3. K. Upadhya, J.-M. Yang, and W. P. Hoffmann, American Ceramics Society Bulletin 76, 51 (1997).  4. A. S. Brown, Aerospace America 35, 20 (1997).  5. S. Norasetthekul, P. T. Eubank, W. L. Bradley, B. Bozkurt, and B. Stucker, Journal of Materials  Science 34, 1261 (1999).  6. A. K. Kuriakose and J. L. Margrave, Journal of  the Electrochemical Society 111, 827 (1964).  7. J. B. Berkowitz-Mattuck, Journal of  the Electrochemical Society 113, 908 (1966).  8. W. C. Tripp, H. H. Davis, and H. C. Graham, American Ceramics Society Bulletin 52, 612 (1973).  9. F. Monteverde and A. Bellosi, Journal of  the Electrochemical Society 150, B552 (2003).  10. W. G. Fahrenholtz, Journal of  the American Ceramics Society 90, 143 (2007).  11. A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, Journal of the European Ceramics Society 27, 2495  (2007).  12. D. Sciti, M. Brach, and A. Bellosi, Scripta Materialia 53, 1297 (2005).  13. D. Sciti, M. Brach, and A. Bellosi, Journal of Materials Research 20, 922 (2005).  14. S. Q. Guo, T. Nishimura, T. Mizuguchi, and Y. Kagawa, Journal of  the European Ceramics Society  28, 1891 (2008).  15. T. Kimura, K. Nishida, and S. Tanuma, Microchimica Acta 155, 175 (2006).  123  \\x0c', 'Oxid Met (2009) 72:335-345  345  16. T. Kimura, T. Sugizaki, K. Nishida, N.  Ishikawa, and S. Tanuma, Journal of  the Japan Institute of  Metals 68, 8 (2004).  17. A. E. McHale, Phase Diagram for Ceramics (American Ceramic Society, Westerville, OH, 1994),  p. 174.  18. P. Holtappels, U. Vogt, and T. Graule, Advanced Engineering Materials 7, 292 (2005).  123  \\x0c']"
},{
  "_id": 231,
  "PDF": "Rapid Oxidation Characterization of Ultra-High Temperature Ceramics.pdf",
  "Text": "['Rapid Oxidation Characterization of Ultra-High Temperature Ceramics  Sigrun N. Karlsdottir  w  and John W. Halloran  Materials Science and Engineering Department, University of Michigan, Ann Arbor, Michigan 48109-2136  Here, a novel method for testing ultra-high-temperature ceram ics (UHTC) at a high temperature, rapidly, at a low cost is in troduced. A self-supported, self-heated ribbon specimen is used  with a table-top apparatus to achieve the necessary high tem peratures. This method enables a large temperature-time-com position parameter space to be covered by rapidly producing a  large set of postoxidation samples  for analysis. The complex  oxide scale known to form during oxidation of UHTC materials  is shown to be easily reproduced using this method. A ZrB2-SiC tested at 17001C for 15 min. (15 vol%) UHTC material was The oxide scale consists of a thin outermost silica (SiO2) layer and a zirconia (ZrO2) columnar layer with small amounts of SiO2 embedded between the ZrO2 columns. A region of SiCdepleted zone was observed between the unreacted core and the ZrO2 layer. The measured thickness of the oxide scale was 102 lm and B120 lm for the SiC-depleted zone.  I.  Introduction  RECENTLY,  interest  in  ultra-high  temperature  ceramics  (UHTC) has increased significantly due to the drive to pro duce a reusable thermal protection system and other components for future generations of hypersonic aerospace vehicles.1  Most modern designs of hypersonic vehicles incorporate sharp  aero-surfaces to increase aerodynamic performance. These de signs  require materials capable of operating in an extreme re entering environment, such as an oxidizing atmosphere at temperatures above 17001C and corrosive gases at high velocities.2  Today,  there are few fully developed materials  that meet  the  needs of thermal protection systems for sharp aero-surfaces and  so UHTC are receiving more attention as the solution.  UHTC are a class of refractory materials including transition  metal borides, carbides, and nitrides. The refractory borides, i.e.  ZrB2 and HfB2, have extremely high melting 430001C, high thermal conductivity, high hardness, and retained strength and chemical stability at elevated temperatures.3  temperatures,  As early as 1953, ZrB2and HfB2-based composites were identified by Hoffman,4 followed by groups led by Berkowitz-Mattuck5 and Kaufman et al.,6 as the most promising material for hypersonic applications. In the 1970s, Kaufman et al.7,8 showed  that by addition of SiC to ZrB2 and HfB2, the oxidation resistance could be greatly improved. Tripp9 and Tripp and Graham10  confirmed Kaufman’s  results and provided additional  information such as microstructural studies of the oxide scale. Past studies by Levine et al.,11 Nguyen et al.,12 Opila et al.,13 Monteverde et al.,14-18 Scatteia et al.,19 Monteverde and Bellosi,20 Fahrenholtz Hilmas and colleagues,21,22 and others23,24  have also shown that addition of SiC to ZrB2 and HfB2 increases densification, thermal shock, and oxidation resistance of the  composites. These researchers have discovered that ZrB2-SiC and HfB2-SiC composites form a complex oxide scale after oxidation at elevated temperatures. The oxide scale is composed of  refractory oxide skeleton and amorphous glass components that  provide oxidation resistance of  the composites at high temper atures.  Despite the  fact  that  it has been almost 50 years  since  re searchers started to explore UHTC materials the oxidation mech anisms of  these materials are still not well understood. This  is  primarily due to the complexity of  the oxide scale and limited  information. Also, much of the oxidation experiments that have temperatures below 17001C. The main  been conducted were at  exception is NASA’s Arc Jet Facility used to simulate aerodynamic heating. It is capable of reaching temperatures from 17001 to 25001C.1,21,25,13 Arc jet  testing is expensive due to the high  power used during testing (40-75 MW), and the facilities are op erated primarily to support the government aerospace research and developmental testing.25 Nongovernmental groups can apply for its use, but it is both time consuming and complicated.26  Here, we report a novel method for testing UHTC materials  at high temperatures rapidly at low cost. A self-supported, self heated ribbon specimen is used with a table-top apparatus  to  achieve the necessary high temperatures  for UHTC oxidation  experiments. The present work introduces the design of the sys tem and function. Also, the complex oxide scale that forms dur ing oxidation of UHTC is shown here to be easily reproduced  using the ribbon method.  II.  The Ribbon Method  (1)  The Self-Supported Ribbon Specimen  The design of the table-top apparatus is based on the fact that  the UHTC materials are metallic conductors. A miniaturized  self-supportive UHTC specimen is resistively heated by passing  a modest current. The specimen is  fabricated by reducing the  thickness of a matchstick-size sample in the center ribbon with a thickness of 400-500 mm. The  to make a  ribbon, herein  called the hot zone, can reach very high temperatures when the  current  is passed into the  thicker  ends of  the  specimen. The  thicker ends of  the specimen will remain relatively cool, while  the hot zone can reach the desired temperature. With this ge ometry the specimen is self-supportive (Fig. 1). The hot zone is  thus not in contact with foreign material. This is important, be cause at  these high temperatures, the specimen may react with  the materials  it contacts. The specimen is  in open air during  testing, and it has a small heat load, and so diagnostics can be  brought close to it during oxidation.  (2)  The Table-Top Apparatus  A table-top apparatus provides  the  current and controls  the  temperature of the specimen. The control unit of the apparatus (MHI BPAN-Ot, Micropyretics  consists of a  control panel  Heaters  Int., Cincinnati, OH), and a programmable tempera ture controller  (Model 2416, Eurotherm, Leesburg, VA). The  power controller is a single-phase silicon (Si)-controlled rectifier  (SCR) with an advanced current limit and a soft start feature. A  stepdown transformer is a part of the apparatus and is connect ed to the controller to give the desired secondary current in the  D. Butt—contributing editor  This work was financially supported by the Office of Naval Research under the grant N  0014-02-1-0034.  Presented at the 108th Annual Meeting of The American Ceramic Society (part of the  MS&T ’06 Materials Science & Technology Conference and Exposition), Cincinnati, OH,  2006.  w  Author to whom correspondence should be addressed. e-mail: nanna@umich.edu  Manuscript No. 22585. Received December 14, 2006; approved May 3, 2007.  Journal  J. Am. Ceram. Soc., 90 [10] 3233 - 3238 (2007)  DOI: 10.1111/j.1551-2916.2007.01861.x  r 2007 The American Ceramic Society  3233  \\x0c', '3234  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 90, No. 10  specimen sitting on a simple (copper) Cu/(silver) Ag bridge fixed  with alligator  clips, which can serve as voltage probes. The  bridge consists of Cu plates that hold up Ag sheets where the  thicker ends of the specimen sit. The Cu plates are connected to  the current  leads and fixed to a thermal  insulation base plate  s  (Transite  HT, Monaco, MI). Ag was used for connecting the  Fig. 1.  Digital image of the self-supported, self-heated ribbon specimen.  range of 0-125 A. The power input used to heat the ribbon specimen to temperatures from 9001 to 20001C is around 90-125  W, where the voltage across the specimen is on the order of 1 V  and the current ranges from 90 to 125 A.  The temperature sensor of the apparatus is a micro optical infrared pyrometer that can measure temperature from 9001 to 33001C with a fast response time (minimum 1 ms). The micro pyrometer is a single-color pyrometer with a wavelength of 1 mm and a bandwidth of 0.7-1.1 mm. The micro pyrometer is focused  on the hot zone of  the specimen and provides the temperature  signal  to the  temperature  controller. The  amount of  current  passed through the sample is controlled with the power control ler to obtain the desired temperature-time schedule. The UHTC  specimen can be quickly heated and cooled due to the fast re sponse of the temperature controller and the pyrometer and the  size of  the miniaturized sample. Any time-temperature profile,  cyclic, or static oxidation can be conducted using the pyrometer  signal as the control variable. With the table-top apparatus and  the ribbon specimen, oxidation experiments at temperatures in the range of 9001-20001C can be performed, without creating  a difficult-to-manage heat  load in the surroundings due to the  small size of the specimen. A similar approach was used by Cabrera and Kirner27 for the design of a cyclic metal-oxidation apparatus. Cabrera and Kirner27 resistively heated thin Si-coated steel foils, 50-mm thick, but these were only capable of reaching temperatures in the range of 4001-10001C.  Figure 2(a) shows the self-supported miniaturized specimen  connected to the  table-top apparatus. Figure 2(b)  shows  the  specimen to the Cu plates due to its high conductivity and ox idation resistance (Cu was not suitable due to its poor oxidation  resistance). Thin alumina (Al2O3) sheets were used between the alligator clips and the sample as heat shields, to avoid over  heating of the alligator clips.  Temperature, current, and voltage are recorded by a LabView  DAQ. An AC clamp on adapter is clamped on the secondary wire  in the circuit to detect the current going through the specimen, and  then transfers it to a data acquisition card (DAQ) (NI USB-6009,  National Instruments Corporation, Austin, TX) connected to a  computer. The temperature measured with the micro pyrometer is  transferred to the computer through a RS-232 cable. To process  the data from the DAQ and the RS-232 cable, LabVIEW from  National  Instruments  is used. A LabVIEW program was de signed to collect and save the temperature, current, and voltage  data sampled during testing. The LabVIEW program samples the  signal from the DAQ and the RS-232 cable (connected to the py rometer) and plots the RMS current versus time and the temper ature versus time. The data are collected at a sampling frequency  of 4 Hz,  i.e. the sampling time is 250 ms. Figure 3 shows an ex ample of a temperature and a current profile of a ribbon specimen tested in the apparatus at 16001C for 15 min.  III.  Experimental Procedure  (1) Material Fabrication  The UHTC material, ZrB2—15 vol% SiC, used in this work was provided by The Institute of Science and Technology for Ce ramics (CNR-ISTEC) in Faenza, Italy. Details of the properties and processing are presented in more detail elsewhere.28 The  fabrication of the self-supporting specimens is performed in two  steps. Firstly, the bulk material is cut with a wire-EDM machine  Fig. 2.  (a) The self-supported ribbon specimen (white box) and its table-top apparatus, (b) the specimen sitting on a copper (Cu)/silver (Ag) holder  (magnification of the white box in the lower left corner of (a)).  \\x0c', 'October 2007  Rapid Oxidation Characterization of Ultra-High Temperature Ceramics  3235  Fig. 3.  (a) Temperature profile, (b) and a current profile of a specimen tested at 16001C for 15 min.  into 2.3 mm \\x02 2.0 mm \\x02 25 mm bars. The bars are then reduced in thickness in the center with a mechanical grinder (220 grit—  diamond wheel) to make the thin ribbon hot thickness of 400-500 mm; see Fig. 1.  zones, with a  (2)  Oxidation Testing  To verify whether the ribbon specimen method could reproduce  the complex oxide scales that form during high-temperature ex periments on UHTC, ZrB2-SiC specimens (CNR-ISTEC, Italy) were tested in the high-temperature apparatus. The specimen was tested at 17001C in stagnant ambient air for 15 min, at a heating rate of 4801C/min and free cooling (7671C/s). The fast  cooling rate is due to the fast thermal response time (2.98 ms at 17001C) of the specimen, which is dependent on the thickness of the specimen (t 5 420 mm) and its thermal diffusivity (1.41 \\x02 10\\x005 m2/s at 17001C). The fast heating minimizes oxidation before the isothermal run at 17001C.  the  The tested specimen was stored in moisture-free desiccators.  A cross section of the oxidized specimen was prepared for mi crostructural analysis by nonaqueous polishing procedures down to a 1 mm finish. The composition and morphology of  the multilayer oxide scale formed after the oxidation test were  characterized by scanning electron microscopy (SEM; Philips  XL30, Hillsboro, OR),  backscattering  electron microscopy  (BSE; Philips), X-ray energy dispersive  spectroscopy (XEDS;  UTW Si-Li Solid State X-ray Detector with integrated EDAX  Phoenix XEDS system, Mahwah, NJ), and electron microprobe  analyzer  (EPMA; Cameca SX100 Microprobe, Gennevilliers,  France). Analyses were performed on the surface of  the speci IV.  Results and Discussion  When monolithic ZrB2(s) oxidizes an oxide scale composed of zirconia, ZrO2(s), and boron oxide liquid, B2O3(l), forms by the reaction16  ZrB2 ðsÞ þ 5=2O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðl Þ  (1)  The liquid B2O3 forms a continuous passive layer acting as a barrier to oxygen (O) diffusion, which results in passive oxida kinetics  below ca.  oxidation  formation) and mass  tion of ZrB2 and a parabolic 10001C. Above 10001C, the oxidation resistance of ZrB2 is poor due to volatilization of B2O3.21 The kinetics is then controlled by the competition between mass gain (ZrO2(s) and B2O3(l) (B2O3(g) vaporization). At higher temperatures, ZrB2 actively oxidizes due to volatilization of the B2O3 liquid, and results in a porous, nonprotective ZrO2 layer.5,16,21-23 Researchers have reported that by adding SiC to ZrB2, the oxidation resistance of the composite above 12001C is improved by the formation of a less volatile silica (SiO2)-rich glass on the exposed surface1,10,21,11,13,18:  loss  SiCðsÞ þ 3=2O2 ðgÞ ! SiO2 ðsÞ þ COðgÞ  (2)  The SiO2-rich layer has been reported to provide passive oxidation resistance to at least 15001C due to the lower volatility of  SiO2 compared with B2O3 at these temperatures. Figure 4 shows an SEM image of the surface of  the tested  ZrB2-15 Fig. 4(a)  vol% SiC specimen. The  bright  peaks  shown  in  indicate ZrO2 peaks embedded in a SiO2-rich glass matrix, verified by XEDS analysis (Fig. 4(b)). A BSE image of  men and cross section.  the cross section of the specimen (Fig. 5) shows that the oxide  Fig. 4.  Surface view of the ZrB2/SiC composite tested at 17001C for 15 min, (a) zirconia (ZrO2) peaks embedded in a silica (SiO2) glass matrix, (b) X-ray energy dispersive spectroscopy analysis of the surface.  \\x0c', '3236  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 90, No. 10  Fig. 5.  Image of a cross section of a ZrB2/SiC composite tested at 17001C for 15 min. Multilayer oxide scale, consisting of an unaffected substrate, a SiC-depleted zone, a columnar zirconia (ZrO2) layer, and a thin superficial layer of silica (SiO2) glass.  scale of the specimen is multilayered and multicomponent. The  chemical composition of  the oxide scale was characterized by  SiO2 glass embedded between the ZrO2 columns. The thickness of the oxide scale was measured to be on average 102 mm. A  XEDS and EPMA techniques. The oxide scale consists of a thin  region of  the SiC-depleted zone was observed between the un superficial  layer of a SiO2 glass containing a small amount of particulate ZrO2; the second layer mainly consists of columnar ZrO2 grains. The columnar ZrO2 layer has a small amount of  reacted core and the ZrO2 layer. The SiC-depleted zone is large ca. 120 mm, which results in a small region of unreacted material  (core). This is evident from the elemental maps from the EPMA  Fig. 6.  (b) O map,  (a) Backscattering electron microscopy (BSE)  image of the cross section of the ZrB2/SiC specimen tested at 17001C for 15 min, showing the oxygen distribution in the silica (SiO2) and zirconia (ZrO2) scales, (c) Zr map, (d) Si map, showing the large SiC-depleted region and the thin strip of the unreacted core, (e) B map, showing the thickness of the oxide scale and the unreacted ZrB2 in the SiC-depleted zone. The white dashed line in (b-e) outlines the thickness of the specimen. The white color in the maps indicates the higher intensity of X-rays detected for each element, while the  black region represents zero detected intensity of the element.  \\x0c', 'analysis (Fig. 6). The existence of a SiC-depleted layer in oxi dized ZrB2/HfB2-SiC composites has been reported by other researchers in the field.10,17,21,25  The EPMA analyses were performed on the  cross  section  shown in the BSE image in Fig. 6(a). The analyses were mainly  performed to verify the results from the XEDS on the distribu tion of the elements: Si, O, zirconium (Zr), and boron (B). Also,  B could not be detected with the SEM/XEDS machine used;  thus, EPMA analyses were performed to give us information on  whether there existed any borosilicate glass or B2O3 in the oxide scale. B2O3 was not detected in the SiO2-rich scale or in the ZrO2 columnar region with EPMA; this is likely due to the small  amount of B2O3 existing in the oxide scale, which could not be detected with the scanning speed (20 ms) used for the EPMA  analysis. The small amount of B2O3 is expected due to the high temperature (17001C) that the specimen experienced. B2O3 is believed to start volatilizing extensively at temperatures above 12001C for a monolithic material.16 Also,  the outermost  layer,  the thin SiO2-rich scale, will not have the same effective role in limiting O diffusion at 17001C as at lower temperatures. The use  of SiO2-forming ceramics at high temperatures has been reported to be limited to a maximum of 17251C because of rapid ox idation, potential volatility, scale melting, and scale/substrate reactions of the SiO2 formed.29 From the O, Zr, and Si maps  (Figs.  6(a)-(e)),  the  small  amount of  interstitial SiO2 glass in the ZrO2 can be easily identified. These maps also show clearly the thin  columnar  layer  outermost SiO2-rich layer (Figs. 6(b)-(d)). As mentioned above, the SiO2 scale contains small amounts of particulate ZrO2. A higher magnification of the SiO2-rich scale and the particulate ZrO2 is shown in Fig. 7; small peaks of the particulate ZrO2 are clearly evident. These peaks were also observed from the surface  of the specimen (Fig. 4(a)). The nature of the particulate ZrO2 peaks and their formation mechanism are discussed elsewhere.30  Our results are consistent with the results reported by other researchers in the field. Levine et al.11 found that the oxide scale of a ZrB2-SiC (20 vol%) composite, tested in air at 19271C for ten 10-min cycles, was composed of large ZrO2 grains in a SiO2rich glassy phase, where the outermost part of the oxide scale  contained less SiO2 glass than the inner most part of the scale. They also found that after oxidation testing in air at 16271C for  ten 10-min cycles (total 100 min),  the oxide scale of  the com posite was composed of  two layers: an outer SiO2 layer and a ZrO21SiO2 layer. A porous SiC-depleted layer was also observed beneath the oxide scale. Gasch et al.25 performed testing  on HfB2-SiC (20 vol%) composites in a simulated re-entry environment using the NASA Ames Arc Jet Facility. The specimens were tested at 16901C for two 10-min cycles. Gash et al.25  concluded that,  for the two conditions, a passive oxidation of  the SiC plays a role in determining the steady-state surface temperature below 17001C. This is consistent with our results;  the  oxide scale of the ZrB2-SiC (15 vol%) specimen tested at 17001C for 15 min consisted mainly of a ZrO2 columnar layer with a small amount of SiO2 embedded, and a thin outermost SiO2 layer. The large SiC-depleted zone observed for the specimen  indicates that the SiO2 layer does not provide enough protection for a passive oxidation at 17001C. This is consistent with the prediction by Jacobson29 about the limited use of SiO2-forming ceramics at higher temperatures (maximum 17251C).  V.  Conclusion  The complex oxide scale known to form during oxidation of  UHTC is  shown to be easily reproduced by using the ribbon  specimen method. The oxide scale of a ZrB2-SiC (15 vol%) specimen tested at 17001C for 15 min consisted of a thin out ermost  SiO2 a ZrO2 amounts of SiO2 embedded between the ZrO2 columns. A region of SiC-depleted zone was observed between the unreacted  layer  and  columnar  layer with  small  core and the ZrO2 layer. The measured thickness of the oxide scale was 102 mm and the B120 mm for SiC-depleted zone.  The table-top apparatus with its ribbon specimen enables us to perform oxidation experiments at temperatures above 9001C and as high as 20001C with power input only around 100 W,  without  creating  a difficult-to-manage heat  load in the  sur roundings due to the small size of  the specimen. This new de sign is a valuable alternative for exposure of UHTC to high  temperatures.  Acknowledgments  The authors would like to thank the following for their contri butions in support of  this work: Dr. Alida Bellosi and her as sociates at CNR-ISTEC in Italy for providing the ZrB2-SiC material, Carl Henderson for support in EPMA analysis, and  Prof. Albert J. Shih and Jia Tao with EDM machining.  References  1M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurface: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 2P. Kolodziej,  ‘‘Aerothermal Performance Constraints for Hypervelocity Small  Radius Unswept Leading Edges and Nosetips’’; NASA Technical Memorandum,  112204, July 1997. 3R. Telle, L. S. Sigl, and K. Takagi,  ‘‘Transition Metal Boride Ceramics’’; pp.  803-945 in Handbook of Ceramic Hard Materials, Vol. 2, Edited by R. Reidel.  Wiley-VCH, Weinheim, Germany, 2000. 4C. A. Hoffman,  ‘‘Preliminary Investigation of Zirconium Boride Caramels  for Gas-Turbine Blade Applicatons’’; NASA Technical Memorandum E52L15a,  Lewis Flight Propulsion Laboratory, Cleveland, OH, 1953. 5J. B. Berkowitz-Mattuck,  ‘‘High-Temperature Oxidation, III. Zirconium and  Hafnium Diborides,’’ J. Electrochem. Soc., 113, 908-14 (1966). 6L. Kaufman, E. V. Clougherty, and J. B. Berkowitz-Mattuck,  ‘‘Oxidation  Characteristics of Hafnium and Zirconium Diboride,’’ Trans. Metall. Soc. AIME,  239 [4] 458-66 (1967). 7L. Kaufman, ‘‘Boride Composites-A New Generation of Nose Cap and Lead ing Edge Materials  for Reusable Lifting Reentry Systems,’’ AIAA Advanced  Space Transportation Meeting. Am. Inst. Aeronaut. Astronaut., 270-8 (1970). 8L. Kaufman and H. Nesor, ‘‘Stability Characterization of Refractory Materials  under High Velocity Atmospheric Flight Conditions, Part I, Vol. I, Summary’’;  Technical Report no. AMFL-TR-69-84, Air Force Materials Laboratory, Wright Patterson Air Base, OH, 1970. 9W. C. Tripp and H. C. Graham,  ‘‘Thermogravimetric Study of Oxidation of  ZrB2 Soc., 118, 1195-971 (1971). 10W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of SiC Addition on the  in Temperature Range of 800 Degrees  to 1500 Degrees,’’ J. Electrochem.  Oxidation of ZrB2,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 11S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem,  ‘‘Evaluation  of Ultra-High Temperature Ceramics  for Aeropropulsion Use,’’  J. Eur. Ceram. Soc., 22, 2757-67 (2002). 12Q. N. Nguyen, E. J. Opila, and R. C. Robinson,  ‘‘Oxidation of Ultrahigh  Temperature Ceramics in Water Vapor,’’ J. Electrochem. Soc., 151 [10] B558-62  (2004). 13E. J. Opila, S. R. Levine, and J. Lorincz, ‘‘Oxidation of ZrB2and HfB2 Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39,  5969-77 (2004).  Fig. 7.  Backscattering electron microscopy (BSE)  image of  the silica  (SiO2) outermost layer containing zirconia (ZrO2) particulate peaks.  October 2007  Rapid Oxidation Characterization of Ultra-High Temperature Ceramics  3237  \\x0c', '3238  Journal of the American Ceramic Society—Karlsdottir and Halloran  Vol. 90, No. 10  14F. Monteverde, A. Bellosi, and S. Guicciardi,  ‘‘Processing and Properties  22W. G. Fahrenholtz,  ‘‘The ZrB2 Volatility Diagram,’’ J. Am. Ceram. Soc., 88  of Zirconium Diboride-Based Composites,’’  J. Eur. Ceram. Soc.,  22,  279-88  (2002). 15F. Monteverde, S. Guicciardi, and A. Bellosi,  [12] 3509-12 (2005). 23M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey,  ‘‘Advances in Microstructure  ‘‘Mechanical, Thermal, and Oxidation Properties of Refractory Hafnium and  and Mechanical Properties of Zirconium Diboride Based Ceramics,’’ Mater. Sci.  Eng., A346, 310-9 (2003). 16F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 17F. Monteverde and A. Bellosi,  ‘‘The Resistance to Oxidation of HfB2-SiC Composite,’’ J. Eur. Ceram. Soc., 25, 1025-31 (2005). 18F. Monteverde, ‘‘The Thermal Stability in Air of Hot Pressed Diboride Matrix  in Dry  Composites for Uses at Ultra-High Temperatures,’’ Corrosion Sci., 47, 2020-33  (2005). 19L. Scatteia, R. Borelli, G. Marino, A. Bellosi, and F. Monteverde,  ‘‘Charac terization and Process of New Metal Diboride Compound for TPS Applications’’;  AIAA/CIRA 13th International Space Planes and Hypersonic Systems and Tech nology Proceedings. American Institute of Aeronautics and Astronautics, Reston,  VA, 2005. 20F. Monteverde and A. Bellosi, ‘‘Development and Characterization of Metal Diboride-Based Composites Toughened with Ultra-Fine SiC Particulates,’’ Solid  State Sci., 7 [5] 622-30 (2005). 21A. Chamberlain, W. Fahrenholtz, G. Hilmas, and D. Ellerby,  ‘‘Oxidation of  ZrB2-SiC Ceramics under Atmospheric and Reentry Conditions,’’ Refractories Appl. Trans., 1 [2] 1-8 (2005).  Zirconium Compounds,’’ J. Eur. Ceram. Soc., 19, 2405-14 (1999). 24C. R. Wang, J. M. Yang, and W. Hoffmann, ‘‘Thermal Stability of Refractory  Carbide/Boride Composites,’’ Mater. Chem. Phys., 74, 272-81 (2002). 25M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties  and Arc  Jet Oxidation  of Hafnium Diboride/Silicon  Carbide Ultra High  Temperature Ceramics,’’  J. Mater.  Sci.,  39,  5925-37  (2004). 26I. Terrazas-Salinas  and C. Cornelison,  ‘‘Test  Planning Guide  for ASF  Facilities’’; A029-9701-XM3 Rev. B Thermophysics Facilities Branch,  Space  Technology Division, NASA Ames Research Center, CA, March, 1999. 27A. L. Cabrera and J. F. Kirner,  ‘‘A Rapid-Temperature-Cycling Apparatus  for Oxidation Testing,’’ Oxidation Metals, 35 [5/6] 471-9 (1991). 28S. N. Karlsdottir, J. W. Halloran, F. Monteverde, and A. Bellosi, ‘‘Oxidation  of ZrB2-SiC: Comparison of Furnace Heated Coupons and Self-Heated Ribbon  Specimens (in 29N.  S.  press).  Jacobson,  ‘‘Corrosion  of  Silicon-Based Ceramics  in Combustion  Environments,’’ J. Am. Ceram. Soc., 76 [1] 3-28 (1993). 30S. N. Karlsdottir, J. W. Halloran, and A. N. Grundy, ‘‘Zirconia Transport by  Liquid Convection during Oxidation of Zirconium Diboride-Silicon Carbide,’’  J. Am. Ceram. Soc., (in press; 2007).  &  \\x0c']"
},{
  "_id": 232,
  "PDF": "REACTIVE HOT PRESSING AND OXIDATION BEHAVIOR OF Hf-BASED ULTRA-HIGH-TEMPERATURE CERAMICS.pdf",
  "Text": "['July 23, 2010  16:52 WSPC/S0218-625X  S0218625X10013886  Surface Review and Letters, Vol. 17, No. 2 (2010) 215-221 c(cid:1) World Scientific Publishing Company  DOI: 10.1142/S0218625X10013886  REACTIVE HOT PRESSING AND OXIDATION  BEHAVIOR OF Hf-BASED ULTRA-HIGH TEMPERATURE CERAMICS  SEUNG JUN LEE∗ , EUL SON KANG† , SEUNG SU BAEK† and DO KYUNG KIM∗,‡ ∗Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology (KAIST), 335 Gwahangno, Yuseong-gu, Daejeon 305-701, Republic of Korea †Agency for Defense Development (ADD), 462 Jochiwongil, Yuseong-gu, Daejeon 305-600, Republic of Korea ‡dkkim@kaist.ac.kr  Received 21 January 2009  A HfB2 -SiC ceramics were  fabricated via a reactive hot pressing using Hf, B4C, and Si as        precursors. The reaction temperature for  the reactive hot pressing between 1800 and 1900  C  was determined by reaction of the precursor at different temperatures from 900 to 1800  C. The  effective size reduction of precursors was  investigated by vibration milling, which exhibited a  critical role to achieve high densification and uniform microstructure. Also,  it has affected the  oxidation behavior of HfB2 -SiC in air. Vibration milled sample showed uniform surface of SiO2  layer, but sample which was fabricated by as-received powder exhibited non-uniform oxidation  behavior. Examination of  the mechanical properties  showed that particle  size  reduction via  vibration also led to improved flexural strength, hardness and fracture toughness.  Keywords : Reactive hot pressing; boride; vibration milling; oxidation.  1.  Introduction  Transition-metal borides have gained considerable importance as ultra-high-temperature ceramics (UHTCs) because of their advantageous characteristic of melting points >3000 C, oxidation resistance and thermal shock resistance.1,2 UHTCs have been extensively investigated for high temperature structural applications that include thermal protection materials for atmospheric re-entry, hypersonic flight, and rocket propulsion. UHTCs materials, such as zirconium diboride (ZrB2) and hafnium diboride (HfB2 ) exposed to air below 1100 C undergo oxidation to HfO2 and B2O3 (l). And the addition of SiC  in matrices of ZrB2 and HfB2 improved the oxidation resistance by forming an SiO2 layer on the surface up to 1700 C.3,4 Transition-metal borides typically need relatively (>1900 C) high temperature and high pressure (>30 MPa) to obtain dense samples because of strong covalent bonding and low diffusion coefficient.2 Therefore, metallic additives have been used to promote liquid-phase formation that can reduce densification temperature, but it can deteriorate properties.5 Recently, the high temperature pressureless sintering of transition-metal borides has been widely investigated by using WC, B4C, C,  ‡  Corresponding author.  215  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  216  S. J. Lee et al.  and MoSi2 .6-8 In contrast to sintering aid assisted sintering methods, spark plasma sintering (SPS) has been reported to possess advantages, such as fast heating rate, short sintering time, high density, and clean grain boundaries.9 Another alternative advantageous method is reactive hot pressing (RHP), which involves both the synthesis and densification into single-step process leading to high density ceramics at reduced temperatures, and lower impurity contents compared with conventional process.10,11 Many reports stated that microstructure of the material fabricated by reactive hot pressing is strongly co-related with particle size of the precursor powder. The present work reports on the fabrication of HfB2 -SiC composite via a RHP technique. The composite fabrication temperature range was investigated by reaction of the precursor at 1800 C. Also, different temperatures from 900 to effects of precursor powder on the densification, oxidation behaviour and mechanical properties were investigated.  2.  Experiment  2.1.  Powder processing  The precursor powder were Hf (purity 99.5%, particle size — 325 mesh, Sigma-Aldrich, USA), B4C (Grade HS, particle size 0.8 µm, H. C. Stack, Germany) and Si (purity 99%, particle size — 325 mesh, Sigma-Aldrich, USA). The chemical reaction can be expressed by the following equation to prepare HfB2 -SiC composite: 2Hf + B4C + Si → 2HfB2 + SiC.  (1)  The volume contents of HfB2 -SiC composite were 74.08 and 25.92 vol.%. The theoretical density of the composite with respect to the rule of mixture is 9.14 g/cm3 . Two milling methods were adopted, in which one to mix and another to reduce the particle size of the precursor powders. In the first method, the precursor powder was ball-milled in ethanol for 24 h using ZrO2 as a milling media and designated as HS. In another method, the precursor powder was vibration milled using ZrO2 as a milling media in a Teflon jar for 2 h to reduce the particle size of the powders and designated as MHS. To investigate the optimum reaction and sintering conditions of the HfB2 -SiC composite with high densification at a low temperature, pressureless heat treatment was conducted.  The powder sample was pressed into disc-shaped pellet followed by cold isostatic pressure (CIP) under 200 MPa and heat-treated in the temperature range of 800-1800 C for 1 h under an Ar atmosphere with a heating rate of 10 C/min. From the above basic experimental condition the composites were sintered using the reactive hot pressing at different temperatures of 1800 and 1900 C for 1 h under pressure of 32 MPa in an Ar atmosphere. The relative bulk density was measured using the Archimedes methods, and phase composition was determined by an X-ray diffractometer (XRD, Rigaku, D/MAX-IIIC X-ray diffractometer, Tokyo, Japan), CuKα radiation (λ = 0.15406 nm at 40 kV and 45 mA). The microstructure of reacted powder mixture and sintered pellets were examined by a scanning electron microscopy (FE-SEM Philips XL30 FEG, Eindhoven, the Netherlands) and a software program(ImageJ). The polished samples were sub jected to the oxidation test under air atmosphere. Each specimen was heat treated at 1500 C for 30 min furnace with a heating rate of 5 C/min. in a tube Hardness and fracture toughness of the samples were determined using Vickers indentation at a load of 2 kg and a dwell time of 15 s. The fracture toughness was estimated through the following equation:  KIC = 0.0319  P  al1/2  where P is the applied load, a the mean indentation half-diagonal length, and l is the crack length. A flexural strength was measured by three point bending method.  3.  Results and Discussion  3.1.  The powder synthesis via  solid-state precursors  Figure 1 shows the XRD patterns of pressureless heat-treated HfB2-SiC composite at different temperatures. The patterns revealed that there was no apparent change in the reaction powder mixture even after heat treated to 800 C for 1 h. The samples heat treated at 1000 C revealed the formation of a cubic phase HfC, in addition with Si and Hf peaks, which indicates the reaction between Hf and B4C occurs at 1000 C. This phenomenon can be explained by conventional reactive hot press model.10,11 At a suitable reaction temperature, it is assumed that B and C  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  Reactive Hot Pressing and Oxidation Behavior of Hf-Based Ultra-High-Temperature Ceramics  217  continues to decrease based on the peak intensity in Fig. 1. At high temperatures, HfC was unstable in the presence of B4C and Si according to reaction (4) and as discussed above, HfC was formed by the reaction between Hf and B4C, and was later consumed in reaction (4):  2HfC + B4C + 3Si → 2HfB2 + 3SiC.  (4)  Reaction (4) is also exothermic and also thermodynamically favorable even at 1000 C. The appearlow intensity XRD peaks at 35.65  indicates ance of the 1600 C, that the formation of SiC occurs near which is higher than HfB2 and HfC because of low reactivity of Si. The XRD pattern of sample reacted at 1800 C showed formation of final composite HfB2 , SiC without any secondary phase. As a result, it can be concluded that the overall powder reaction takes place in two steps, in which the initial reaction initiated the formation of HfC and at higher temperatures, simultaneous reactions take place by reactions (3) and (4) and formed HfB2 , SiC and elimination of HfC. Similar steps were followed for the vibration milled powder MHS to obtain HfB2 -SiC composite. During the milling process 1.5 wt.% of ZrO2 impurity was introduced because of wear of the ZrO2 balls. ZrO2 impurity can be removed by B4C following reaction (5):  5B4C + 7ZrO2 → 7ZrB2 + 5CO (g) + 3B2O3 (l).  (5)  Reaction (5) is well defined in the previous report.6 Reaction between B4C and ZrO2 was initiated around 1200 C and complete reaction occurred at 1450 C to form ZrB2 . In literature, there are reports on the mutual continuous solid solution forms among Groups IV and V of dibiride,12 i.e. ZrB2 could form continuous solid solution with HfB2 . Thus, ZrO2 contamination from ZrO2 balls was accommodated in the HfB2 lattice by the formation of solid solution via additional B4C at the stoichiometric of reaction (5). In case of MHS, it was observed from the XRD pattern that there was negligible peak shift (not shown here); this may be due to a small amount of ZrO2 contaminations. The reaction sequence was similar with HS except for the initiation temperature of reaction between Hf and B4C and formation of HfB2 were and 1000 C, decreased to 800 respectively. Particle size reduction via vibration milling provides the  Fig. 1.  The XRD patterns of the Hf/B4C/Si mixed pow ders which are heat-treated at different temperatures for  60 min in Ar.  atoms from B4C diffuse faster than Hf and Si. Among this the diffusion rate of C is higher than that of B, and easily reacts with Hf as Si possesses lower chemical reactivity than Hf. As a result, it is evident that formation of HfC occurs at lower temperature 1000 C and faster than HfB2 and SiC. Thus, it can be concluded that HfB2 formation required a higher temperature compared with HfC. The formation of HfB2 occurs at a relatively higher temperature above 1200 C by a reaction between Hf and B or residual B4C due to the increment of the diffusion rate of B. By analyzing the reaction sequence, it can be evidently concluded that the reactions (2) and (3) may occur in two steps as follows: Hf + B4C → HfC + 4B 2Hf + 4B → 2HfB2 .  (2)  (3)  The above reactions (2) and (3) possess negative Gibbs free energies of ∆G1000 = −126 and ∆G1000 = −630 kJ/mol, respectively, indicating that the reactions are all thermodynamically favorable. In the XRD patterns, it is observed that the sample heat treated at 1200 C revealed the formation of HfO2 phase. The presence of HfO2 might be attributed to oxide impurities in the precursor powder and/or oxygen uptake during milling process. Further increase in reaction temperature result in the formation of HfB2 as the main phase and the intensity of HfC  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  218  S. J. Lee et al.  shorter diffusion path and this might decrease reaction temperature.  the  3.2. Microstructure  SEM analysis was conducted to investigate particle size and morphology of HS and MHS after the  completion of the reaction. Figure 2(a) shows the fracture surface of HS after reaction at 1800 C, and  Fig. 3(a) shows the microstructure of HS densified at 1900 C for 60 min. After the reaction at 1800 C, HS  consists of  large particles up to 40 µm that contains  sub-micron meter grains and randomly distributed  Fig. 2.  SEM images of the fracture after heat-treated at 1800     C for 60 min (a) HS, and (b) MHS.  (a)  (b)  (a)  (b)  Fig. 3.  SEM micrographs of polished surface and area of SiC phase of sample HS (a), (b) and sample MHS (c), (d).  (c)  (d)  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  Reactive Hot Pressing and Oxidation Behavior of Hf-Based Ultra-High-Temperature Ceramics  219  particles. This result can be correlated with the chemical reaction at 1800 C yield HfB2 and SiC and the large grains might be HfB2 , and the others are HfB2 and SiC. As discussed above, during the reaction process, the diffusion rates of Hf and Si atoms are slow and the compound formation is mainly controlled based on the reactivity and diffusion rates of boron and carbon. Due to the large particle size of the Hf precursor powder and the resultant product HfB2 also formed with similar size grains, in the RHPed composite as shown in Fig. 3(a). Thus, as shown in Fig. 3(a), microstructure is inhomogeneous with large HfB2 (light phase) grains and non-uniform distribution of SiC (dark phase). Size and morphology of formed HfB2 and SiC are related to its original particle size and shape. Although as shown in Fig. 2(a) inset, primary grain of formed HfB2 is sub-micron meter size, the sizes of the secondary HfB2 grains are larger. This large grain might be the formation of residual porosity and, therefore, limit the final densities (96.3% R.D.). The size and 1800 C for morphology of MHS after reaction at 1 h is shown in Fig. 2(b). It can be observed in the Fig. 2(b), the particle size of formed HfB2 is about 5 µm and also consists of sub-micron size primary grains and about 3 µm SiC. The polished surface of RHPed MHS is shown in Fig. 3(c). The pellets sintered via RHP at 1900 C revealed complete densification (99.8% R.D.) with rare residual pores and homogeneous microstructure with uniform distribution of SiC. As shown in Fig. 3(d), average area of the SiC particulate is also reduced in MHS (4.6 µm2 )  compared with HS (20.6 µm2 ). And this area size reduction of SiC also means microstructure of MHS is more homogeneous than HS with well distributed SiC in HfB2 matrix. From this analysis, it can be concluded that the particle size reduction via vibration milling can obtain small size HfB2 compared with as-received powder. Increased driving force due to the formation of small size of HfB2 and application of external pressure during the RHP process led to the nearly high densification without residual pore. In addition to the formation of small particles, increase in the defect concentration owing to the high milling also could possibly be a reason for the increased densification behaviors similar to the study of particle size effects on the sintering of boride compound.13  3.3.  Oxidation behavior  Figure 4 shows the cross-sectional analysis of the oxidized HS and MHS samples heated in air at 1500 C for 30 min. The oxidation behaviors of the two samples appear to be similar, except for the distribution of surface SiO2 layers. Oxidation of MHS produced a structure that consisted of three layers. As shown Fig. 4(b), oxidized cross-section con(3 µm), SiC-depleted layer (10 µm) sists of SiO2 and un-reacted HfB2 -SiC layer. The formation of layered structure is consistent with observations of other investigators who have studied ZrB2-SiC and HfB2-SiC.3,4,14 But in case of HS, distribution of the SiO2 layer on the surface was non-uniform over the  (a)  (b)  Fig. 4.  SEM images of (a) HS showing non-uniform SiO2 distribution, and (b) MHS showing a uniform layer of SiO2  (I), SiC depleted layer (II), and un-reacted HfB2 -SiC layer (III) after exposure of samples at 1500  C for 30 min in air.     Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  220  S. J. Lee et al.  specimen surface. This might be due to the inhomogeneous microstructure which is originated from large starting Hf powder. In the HfB2 rich region, there was no Si source. Thus, monolithic HfB2 was exposed to the 1500 C in air. As HfB2 is exposed to 1500 C in air, it can be possibly oxidized as following the reaction (6)  HfB2 + 5/2O2 → HfO2 + B2O3 (l).  (6)  In literature, there are many reports that describe the high vapour pressure of B2O3 at 1500 C,14 and thus B2O3 will evaporate rapidly leaving a porous HfO2 . If B2O3 evaporates during the oxidation, as is the case for monolithic ZrB2 (or HfB2 ), the effective diffusion barrier is reduced since the porous HfO2 layer alone does not protect the underlying HfB2 from rapid oxidation. In this region, HfB2 oxidation will exhibit linear kinetics, i.e. oxidation depth is increased linearly with time and reaction ratecontrolled kinetics can be predicted. In contrast to locally SiC rich region, parabolic oxidation behavior like MHS was predicted by encouraging the formation of Si rich glass layer on the SiC rich region, as shown in Fig. 4(a). Nonuniform distribution of SiO2 may be due to wetting characteristics and/or local variation such as microstructure with small effective area due to the large SiC area (20.6 µm2 ) that might enhance the local oxidation rate.  3.4. Mechanical properties  The relative density and mechanical properties of the composites are listed in Table 1. The variation in the flexural strength was well reflected in the change of relative density and microstructure features. The incensement of flexural strength in MHS compared with HS due to the finer raw powder via vibration  Table 1.  Relative density and mechanical properties of  the RHPed samples.  and thus, higher density and lower pore distribution with well distribution of SiC phase in the bulks. The fracture toughness of MHS is slightly high compared with the HS. But fracture toughness of HS and HMS is high with the other HfB2 -SiC composite densified by hot pressing.15 The Vickers hardness was 17.7, 17.5 and 18.3 GPa for HS1900, MHS1800 and MHS1900, respectively, showing they were in a similar level. But hardness is increased to 18.3 GPa for MHS due to the full density with no porosity compared with HS1800. Hardness has been shown to decrease exponentially as the porosity increases for ceramic materials.16  4.  Conclusions  This work reported reactive powder synthesis via solid-state precursors, microstructure, oxidation behavior and mechanical properties prepared from a mixture of Hf/B4C/Si powders. The reactions of the powder mixture commenced at 1000 C and completed at 1800 C. A relative density densified with vibration milled power mixture was 99.8% at 1900 C. The enhanced densification and uniform microstructure was attributed formation of fine HfB2 and SiC phases. The size reduction of starting powder gave the homogeneous microstructure with uniform distribution and increased effective area of SiC which play a critical role of improving the oxidation resistance and provided uniform oxidation behavior over the whole specimen. The size reduction of starting powder also led to higher mechanical properties. The mechanical properties for the fully dense RHP materials were comparable to reported value.  Acknowledgments  This work was supported by the Agency for Defense Development under the contract UE075126GD. S. J. Lee would like to thank the National Research Foundation of Korea (NRF) grant funded by the Korea government (MEST) (No. 2008-0062204).  Relative Flexural  Fracture  density  strength  Material  (% TD)  (MPa)       HS 1800  HS 1900  MHS 1800  MHS 1900  C  C       89.5  96.3  97.2  99.8  C  C  297 ± 20 551 ± 40 578 ± 23 692 ± 58  4.17 ± 0.38 4.78 ± 0.42 4.87 ± 0.20 5.23 ± 0.17  14.2 ± 0.4 17.7 ± 0.5 17.5 ± 0.5 18.3 ± 0.3  Vickers  hardness  References  toughness  1/2  (MPam  )  (GPa)  1. A. K. Kuriakose and J. L. Margrave, J. Electrochem. Soc. 111 (1964) 827. 2. K. Upadhya, J.-M. Yang and W. P. Hoffman, Am. Ceram. Bul l. 58 (1997) 51. 3. A. Rezaie, W. G. Fahrenholtz and G. E. Hillmas, J. Eur. Ceram. Soc. 27 (2007) 2495.  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c', 'July 23, 2010 S0218625X10013886  16:52 WSPC/S0218-625X  Reactive Hot Pressing and Oxidation Behavior of Hf-Based Ultra-High-Temperature Ceramics  221  4. P. Lespade, N. Richet and P. Goursat, Acta Astronautica 60 (2007) 858.  5. D. M. Van Wie, D. G. Drewry Jr., D. E. King and C. M. Hudson, J. Mater. Sci. 39 (2004) 5915.  6.  7.  8.  S. C. Zhang, G. E. Hilmas and W. G. Fahrenholtz, J. Am. Ceram. Soc. 89 (2006) 1544.  S. C. Zhang, G. E. Hilmas and W. G. Fahrenholtz, J. Am. Ceram. Soc. 91 (2008) 26. L. Silvestroni and D. Sciti, Scripta Mater. 57 (2007)  165. 9. A. Bellosi, F. Monteverde and D. Sciti, Int. J. Appl. Ceram. Technol. 3 (2006) 24.  10. G.-J. Zhang, Z.-Y. Deng, N. Kondo, J.-F. Yang and T. Ohji, J. Am. Ceram. Soc. 83 (2000) 2330.  11. W.-W. Wu, G.-J. Zhang, Y.-M. Kan and P.-L. Wang, J. Am. Ceram. Soc. 89 (2006) 2967. 12. M. J. Gasch, D. T. Ellerby and S. M. Johnson, Handbook of Ceramics Composites  (Springer, Chap.  9,  2005), p. 197. 13. B. Cech, P. Oliverus and J. Sejbal, Powder Metal l. 8 (1965) 142. 14. W. G. Fahrenholtz, J. Am. Ceram. Soc. 88 (2005)  3509. F. Monteverde, J. Al loys Compd. 428 (2007) 197.  15.  16. V. Milman, S.  I. Chugnova,  Chudoba, W. Lo jkowski and W. Gooch, Refract. Metal Hard Mater. 17 (1999) 361.  I. V. Goncharova, T. Int. J.  Surf. Rev. Lett. 2010.17:215-221. Downloaded from www.worldscientific.comby UNIVERSITY OF NEW ENGLAND LIBRARIES on 10/28/14. For personal use only.\\x0c']"
},{
  "_id": 233,
  "PDF": "Reactive Hot Pressing of ZrB2–SiC–ZrC Ultra High-Temperature Ceramics at 1800 C.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 89 [9] 2967 - 2969 (2006)  DOI: 10.1111/j.1551-2916.2006.01145.x  r 2006 The American Ceramic Society  Reactive Hot Pressing of ZrB2-SiC-ZrC Ultra High-Temperature Ceramics at 18001C  Wen-Wen Wu  z  , Guo-Jun Zhang,*,w  Yan-Mei Kan, and Pei-Ling Wang  State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics,  Shanghai 200050, China  A ZrB2-SiC-ZrC composite was prepared from a mixture of zirconium, silicon, and B4C via reactive hot pressing at a relatively low temperature (18001C) for 60 min under 20 MPa in an argon atmosphere. The relative density was 96.8%, the  micro-hardness (Hv10) was 16.7 GPa, and the fracture toughness was 5.1 MPa . m1/2. The presence of ZrC was helpful the densification process and improved the mechanical proper for  ties of  the composite. A model of  the microstructure develop ment  of  the  composite was  proposed  to  explain  the  phase  distribution.  I.  Introduction  Z IRCONIUM diboride (ZrB2) and hafnium diboride (HfB2) are ultra high-temperature ceramics (UHTCs) that have a number of unique properties, including extremely high melting  temperature and hardness, low volatility, and high thermal and electrical conductivity.1 Studies have shown that composites of  diborides with silicon carbide  (ZrB2/SiC and HfB2/SiC) candidates for ultra high-temperature applications, because the the oxidation resistance.2,3 The  are  addition of SiC can improve  presence of SiC can also increase the strength of by acting as a grain-growth inhibitor.4  the materials  For most of  the reported studies, ceramic composites have  been  fabricated  simply  by  hot  pressing  from commercially  available powders. Reactive hot pressing (RHP) is an alternative  route.  It  can  produce materials with  novel  and  controlled  microstructures, with  high  chemical  compatibility  of  the  in  situ-formed individual phases, and phase distribution uniformity.5 In boride-containing composites, for example, TiB2-SiC binary composites and TiB2-Ti(C,N)-SiC ternary composites have been prepared from a mixture of TiH2, Si, and B4C.6-8 For zirconium boride-containing UHTCs, a high-strength ZrB2-SiC composite has been prepared by RHP at 19001C for 60 min under 30 MPa pressure from a mixture of Zr, Si, and to reaction (1).9 The  B4C according a facture toughness of 4.0 MPa \\x01 m1/2. composite is 97.67% with a Vickers hardness of 21.0 GPa and  relative density of  this  2Zr þ Si þ B4C ¼ 2ZrB2 þ SiC  (1)  The addition of ZrC to ZrB2-SiC to form a ternary composite of ZrB2-SiC-ZrC can tailor the microstructure and properties of ZrB2-SiC, especially the superior resistance to oblation at a high temperature.10 In the current work, a composite in the  ZrB2-SiC-ZrC system was prepared by RHP, using Zr, Si, and B4C as starting powders according to reaction (2):  ð2 þ xÞZr þ ð1 \\x00 xÞSi þ B4C ¼ 2ZrB2 þ ð1 \\x00 xÞSiC þ xZrC  (2)  When x 5 0,  reaction  (2)  reduces  to reaction  (1).  In the  present work, we  took  x 5 0.20  based  on  calculations  for  maintaining  the  content of SiC at  20  vol%. The  calculated  volumetric  composition  is  (vol%)  73.6ZrB2120SiC16.4ZrC. The theoretical density of the composite calculated according to the rule of mixtures is 5.55 g/cm3, based on the densities of 6.09, 3.21, and 6.44 g/cm3 for ZrB2, SiC, and ZrC, respectively. The true density should be lower than the calculated one due to  the impurities (mainly Ti). For comparison, a ZrB2-SiC composite without ZrC (when x 5 0) was also prepared. Reactions  (1) and (2) are thermodynamically favorable and exothermic. In  this communication, the samples were produced under relatively mild conditions (18001C), and the manufacturing process, me chanical properties, and microstructure of  the composites are  reported, and the effect of ZrC on sintering is discussed.  II.  Experimental Procedure  The starting powders were Zr (purity 95.82%, impurities include  Ti 2.34, Hf 0.52, Fe 0.24, W 0.08, Cr 0.06, Ni 0.03, O 1.09, particle size o25 mm, Guoyao Chemicals Co. Ltd., Shanghai, China), Si (purity499%, particle size o50 mm, Yinfeng Silcon Co. Ltd., Jinan, China), and B4C (purity 99%, particle size about 2 mm, Jingangzuan Boron Carbide Co., Ltd, Mudanjiang,  China). The stoichiometric powders were mixed and ground in  ethanol  in an agate mortar until all of  the ethanol was evapo rated. The  grinding  process was  repeated  once  for  a  total  grinding time of about 1 h. The mixed powder was then dried  and placed in a graphite die with a BN coating. The composite was then RHPed at a temperature of 18001C for 60 min under a  pressure of 20 MPa in an argon atmosphere. A slow heating rate (101C/min) was adopted to prevent the reaction from becoming  self-sustaining. The application of pressure was initiated at 15501C. The reacted disk had dimensions of 20 mm (diameter)  and 5 mm (thickness).  After removing the surface layer from the hot-pressed disk by  grinding, the bulk density was measured using the Archimedes  method. Phase composition was determined by X-ray diffractometry (XRD) using CuKa radiation. The disk was ground with  SiC abrasives and then polished using a diamond pastor of 1 mm. The hardness and fracture toughness were measured by  the indentation method, using a load of 10 kg for 10 s on a  2967  W. Fahrenholtz—contributing editor  This work was financially supported from the Chinese Academy of Sciences under the  Program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program),  the  National Natural Science Foundation of China, and the State Key Laboratory of High  Performance Ceramics and Superfine Microstructures of Shanghai Institute of Ceramics.  *Member, American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: gjzhang@mail.sic.ac.cn z Graduate School of the Chinese Academy of Sciences, Shanghai, China.  Manuscript No. 21572. Received March 10, 2006; approved April 26, 2006.  \\x0c', 'polished surface;  the  reported value was  an average of five  measurements.  Scanning  electron microscopy  (SEM)  was  performed to observe the microstructure of the composite.  III.  Results and Discussion  According to the XRD patterns shown in Fig. 1, phases present  in the composites agree with those predicted from reactions (1)  and (2). This observation means that either ZrB2-SiC or ZrB2- SiC-ZrC can be produced by in situ reaction.  The experimental values of some properties of both compo sites are compared in Table I. The relative density (96.8%) of the  ZrB2-SiC-ZrC composite was higher than that of the ZrC-free composite. This indicates that the presence of ZrC improves the  densification of  the material, even though the reactants of  the  two systems are similar.  Compared with the hardness of 21.0 GPa of the composite we mentioned above,9 the micro-hardness (Hv10) of the composite  is not high as listed in Table I. The main reasons for this may be  the  inhomogeneous microstructure  and  the  porosity  in  the  composites. The regions that are rich in SiC will have a higher  hardness than the other regions according to a similar system of TiB2-SiC-TiC.12 The fracture toughness of the composite ZrB2- SiC-ZrC is 5.1 MPa \\x01 m1/2, which was higher than that measured for ZrB2-SiC (4.5 MPa \\x01 m1/2); the formation of ZrC as the third phase improved the fracture toughness of the composite. It is the same case in the TiB2-SiC-TiC system.12 The authors of study concluded that toughness was lower for ZrB2-SiC because of little interaction between the phases, but the presence of ZrC  this  induces strong bonds with both ZrB2 and SiC. Figure 2 shows the XRD pattern of the intermediate products treated at 8001C for 60 min  in the ZrB2-ZrC-SiC system heat in an argon atmosphere. It can be found that ZrC formed at a 8001C. As  relatively  low temperature  of  the  temperature  increased, ZrB2 appeared and the  became  the main  phase;  in  addition,  SiC  amount  of ZrC decreased, based on the  peak intensity in the XRD patterns. By analyzing the reaction  process,6,11  reactions  (1) and (2) may take place  in steps as  follows:  3Zr þ B4C ¼ 2ZrB2 þ ZrC  (3)  2ZrC þ B4C þ 3Si ¼ 2ZrB2 þ 3SiC  (4)  The above reactions have large negative Gibbs free energies (DG298 (3) 5 \\x00767.641, DG1800 (3) 5 \\x00703.078, DG298 \\x00395.212, DG1800 (4) 5 \\x00320.257 kJ/mol), showing that reactions are all thermodynamically favorable.  (4) 5  the  The microstructure of ZrB2-SiC-ZrC composite is shown in Fig. 3. This picture indicates that the distributions of the in situ formed ZrB2, SiC, and ZrC phases homogenous. The particle size of SiC and ZrC is generally small, about 1-3 mm, whereas that of ZrB2 is large, which is about 3-10 mm. SiC and ZrC are mainly located at the  in the composite are not  similar and  ZrB2/ZrB2 grain boundaries, but the phases do not appear in the same region. Namely, regions rich in SiC are poor in ZrC. This  supports the conclusion that the SiC phase is formed from ZrC.  Figure 4 shows a schematic of a possible formation sequence  during RHP. For the model, it is assumed that B and C atoms from B4C diffuse faster than Zr and Si.9,13 Also, Si has a lower reactivity than Zr, so ZrB2 and ZrC formed before any Si compounds. At higher temperatures, SiC formed in situ as a  result of  the reaction between Si, ZrC, and the residual B4C. Owing to the large particle sizes of the Zr starting powder and  its agglomeration, the size of the resulting ZrB2 grains are large  20  30  40  50  60  70  SiC  2 theta (degree)  ZrC  (a)  (b)  ZrB2  Fig. 1.  X-ray diffractometry patterns of the reactive hot-pressed com posites: (a) ZrB2-SiC-ZrC; (b) ZrB2-SiC.  Table I.  Characteristics of the Obtained Composites  Property  Density  (g/cm3)  Relative  density  (%TD)  Hardness  (GPa)  Fracture  toughness (MPa \\x01 m1/2)  ZrB2-SiC ZrB2-SiC-ZrC  5.09  94.8  13.8  4.5  5.37  96.8  16.7  5.1  20  30  40 50 2 theta (degree)  60  70  ZrB2 ZrC  Zr  Si  Fig. 2. X-ray diffractometry pattern of the mixed powder for preparing a ZrB2-SiC-ZrC composite after heat treatment at 8001C for 60 min in an argon atmosphere.  Fig. 3.  Backscattered electron images  of  a  polished  surface  of  the  reactive hot-pressed composite. The gray phase is ZrB2, the dark phase is SiC, and the white phase is ZrC.  2968  Communications of the American Ceramic Society  Vol. 89, No. 9  \\x0c', 'September 2006  Communications of the American Ceramic Society  2969  The microstructure of  the  composite was not homogeneous  owning to the coarse particle size of the starting powders. The  reaction took place in steps. First, B4C reacted with Zr to form ZrB2 and ZrC at a very low temperature. At a higher temperature, SiC was formed in situ by the reaction of Si with ZrC and  the residual B4C.  References  1K. Upadhya,  J.-M. Yang,  and W. P. Hoffman,  ‘‘Materials  for Ultrahigh  Temperature Structural Applications,’’ Am. Ceram. Soc. Bull., 58 [12] 51-6 (1997). 2M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties  and Arc  Jet Oxidation  of Hafnium Diboride/Silicon  Carbide Ultra High Temperature Ceramics,’’ J. Mater. Sci., 39, 5925-37 (2004). 3W. C. Tripp, H. H. Davis, and H. C. Graham,  ‘‘Effect of an SiC Addition on  the Oxidation of ZrB2,’’ Am. Ceram. Soc. Bull., 52 [8] 612-6 (1973). 4A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  ‘‘High-Strength Zirconium Diboride-Based Ceramics,’’ J. Am. Ceram. Soc., 87 [6]  1170-2 (2004). 5M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey,  ‘‘Mechanical, Thermal,  and Oxidation Properties of Hafnium and Zirconium  J. Zhang, Z. Z.  Jin, and X. M. Yue,  Compounds,’’ J. Eur. Ceram. Soc., 19 [13-14] 2405-14 (1999). 6G.  ‘‘Reaction Synthesis of TiB2-SiC Composites from TiB2-Si-B4C,’’ Mater. Lett., 25, 97-100 (1995). 7G. J. Zhang, X. M. Yue, Z. Z. Jin, and J. Y. Dai, ‘‘In-Situ Synthesized TiB2 Toughened SiC,’’ J. Eur. Ceram. Soc., 16, 409-12 (1996). 8G.  ‘‘TiB2-Ti(C,N)-SiC Composites Prepared by Reactive Hot Pressing,’’ J. Mater. Sci. Lett., 15, 26-8 (1996). 9G.-J. Zhang, Z.-Y. Deng, N. Kondo, J.-F. Yang, and T. Ohji,  and Y. M. Yue,  J. Zhang, Z. Z.  ‘‘Reactive Hot  Jin,  Pressing of ZrB2-SiC Composites,’’ J. Am. Ceram. Soc., 83 [9] 2330-2 (2000). 10B. J. White and M. J. Kauffman. U.S. Patent No. 5750450. 11F. Monteverde,  in the Fabrication of Ultra-High-Temperature  ‘‘Progress  Fig. 4. Microstructure  composite  in the  formation mechanism of  the ZrB2-ZrC-SiC reaction-synthesis process, depicting the conversion  from (a)  the powder compact to (b)  the intermediate state, and (c)  the  final microstructure of the composite.  in the RHPed composite. In addition, the regions where SiC is  concentrated are about the same size as the starting Si powders.  Accordingly, using starting powders with a finer particle size and  improving the uniformity of mixing should promote formation  of a more homogeneous material.  From Fig. 3,  by  SiC are  it can be seen that the ZrB2 grains surrounded finer than those not surrounded by SiC. This  supports the conclusion that SiC can reduce grain growth in ZrB2.3,14 This will be useful tions, as a grain-growth inhibitor would improve the material’s  in ultra high-temperature applica stability at  elevated temperature. On the other hand,  the  ex istence  the ZrB2 platelets, and this should have improved the fracture toughness  of ZrC seems  promote  growth  the  to  of  of the composite.  IV.  Summary  A ZrB2-SiC-ZrC composite was prepared by the RHP of a mixture of zirconium, silicon, and B4C powders. The hardness and the fracture toughness of the product measured were 16.7 GPa and 5.1 MPa \\x01 m1/2. The presence of ZrC improves densification process and fracture toughness of the composite.  the  Ceramics:  ‘In Situ’ Synthesis, Microstructure and Properties of a Reactive Hot pressed HfB2-SiC Composite,’’ Compos. Sci. Technol., 65, 1869-79 (2005). 12F. Mestral and F. Thevenot, ‘‘Boride-Carbide Composites: TiB2-TiC-SiC’’; pp. P457-81 in The Physics and Chemistry of Carbides, Nitrides and Borides, Edited  by R. Freer. Kluwer Academic Publishers, Dordrecht, MA, 1989. 13W. G. Fahrenholtz, ‘‘Reactive Processing in Ceramic-Based Systems,’’ Int. J.  Appl. Ceram. Technol., 3 [1] 1-12 (2006). 14R. Telle, L. S. Sigl, and K. Takagi,  ‘‘Boride-Based Hard Materials’’; pp. 803-  945 in Handbook of Ceramic Hard Materials, Edited by R. Ridrel. Wiley-VCH,  Weinheim, Germany, 2000.  &  \\x0c']"
},{
  "_id": 234,
  "PDF": "Refractory Diborides of Zirconium and Hafnium.pdf",
  "Text": "['Refractory Diborides of Zirconium and Hafnium  William G. Fahrenholtz*,w  and Gregory E. Hilmas*  Materials Science and Engineering Department, University of Missouri-Rolla, Rolla, Missouri 65409  Inna G. Talmy* and James A. Zaykoski*  Naval Surface Warfare Center, Carderock Division, West Bethesda, Maryland 20817  This paper  reviews  the crystal chemistry,  synthesis, densifica tion, microstructure, mechanical properties, and oxidation be(ZrB2) refractory diborides  havior  of  zirconium diboride  and  hafnium diboride  (HfB2) complete  ceramics. The  exhibit partial or  solid solution with other  transition metal diborides,  which allows compositional tailoring of properties such as ther mal expansion coefficient and hardness. Carbothermal reduction  is  the typical  synthesis  route, but  reactive processes,  solution  methods, and pre-ceramic polymers can also be used. Typically,  diborides are densified by hot pressing, but recently solid state  and liquid phase  sintering routes have been developed. Fine grained ZrB2 and HfB2 have strengths of a few hundred MPa, which can increase to over 1 GPa with the addition of SiC. Pure  diborides exhibit parabolic oxidation kinetics at below 11001C, but B2O3 volatility leads to rapid, tion kinetics above that temperature. The addition of silica scale  temperatures  linear oxida formers such as SiC or MoSi2 improves the oxidation behavior above 11001C. Based on their unique combination of properties, ZrB2 and HfB2 ceramics are candidates for use in the extreme environments associated with hypersonic flight, atmospheric re entry, and rocket propulsion.  I.  Introduction  Z IRCONIUM diboride (ZrB2) and hafnium diboride (HfB2) are members of a family of materials known as ultra high-temperature ceramics (UHTCs). Several carbides and nitrides of the  group IVB and VB transition metals are also considered UHTCs based on melting temperatures in excess of 30001C and other  properties. Very few elements or compounds from any class of  ceramic materials have melting temperatures 30001C (Fig. 1).1 From the broader  approaching  family of UHTCs,  this  paper focuses on ZrB2 and HfB2. Discussion of TiB2 has been intentionally minimized due to its more widespread use as armor and cutting tools.2-4  Some  structural,  physical,  transport,  and  thermodynamic  properties of ZrB2 and HfB2 are Applications that take advantage of these properties include refractory linings,13-15 electrodes,16-18 microelectronics,19 and cutting tools.20 In addition to high melting temperatures, ZrB2 and HfB2 have a unique combination of chemical stability, high electrical and thermal conductivities, and resistance to erosion/  summarized in Table  I.5-12  corrosion that makes them suitable for the extreme chemical and  thermal environments associated with hypersonic flight, atmospheric re-entry, and rocket propulsion.21-24 Because of  recent  efforts to develop hypersonic aerospace vehicles and re-usable  atmospheric re-entry vehicles,  interest in UHTCs has increased  significantly in the past  few years. As a result, groups  in the  United States, Italy, Japan, India, and China are investigating  diborides. Many of these studies draw inspiration from the late  Professor G. V. Samsonov, whose work is noted for both its  high quality and enormous quantity.  This article provides a critical evaluation of historic (1970s or  earlier) and recent (1990 or later) studies of ZrB2 and HfB2 ceramics as well as presenting some recent results. Fundamental  aspects of crystal structure and bonding, synthesis, densification,  microstructures, properties, and oxidation behavior are exam ined.  II.  Crystal Structure and Bonding  Crystal chemistry and crystal structure determine the chemical,  physical, and thermal properties of materials. For the transition  metal diborides,  fundamental aspects of crystal  structure and  crystal chemistry have been discussed extensively dating to the early 1950s.25-46 This  section reviews  the bonding,  structure,  structure-property relations, and solid solutions of the diborides.  (1)  Bonding and Structure  Borides have a wide range of compositions with boron:metal  (B:M)  ratios  ranging from 1:4 to 12:1. The B:M ratio affects  both properties and electronic structure. Changing B:M changes  the electronic structure of boron, which leads to the formation  of different  structural  complexes  containing one-,  two-,  and  three-dimensional  (3D) B networks.  Increasing the number of  Feature  D. Green—contributing editor  At UMR, portions of  this work were  funded by the Air Force Office of Scientific  Research (F49620-03-1-0072 and FA9550-06-1-0125),  the National Science Foundation  (DMR-0346800),  and  the Air  Force Research  Laboratory  (FA8650-04-C-5704). At  NSWCCD the work was  funded by the Office of Naval Research on several contracts  monitored by Dr. Steve Fishman.  *Member, American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: billf@umr.edu  Manuscript No. 22451. Received November 6, 2006; approved January 10, 2007.  Journal  J. Am. Ceram. Soc., 90 [5] 1347 - 1364 (2007)  DOI: 10.1111/j.1551-2916.2007.01583.x  r 2007 The American Ceramic Society  \\x0c', 'B atoms in a structural complex leads to increases in the B-B  bond strength and an increase in the stiffness of the crystal  lat tice along with increases in melting temperature (Tm), hardness (HV), strength (s), and chemical stability. The M-B bond strength in diborides depends on the degree of  electron localization around the M atoms. The valence electron configuration in isolated B atoms is 2s22p. In metal borides, the outer electron configurations are sp2 and sp3, which promote  strong covalent bonding. In diborides, B atoms are electron ac ceptors, while the M atoms are electron donors. Each M atom  donates two electrons (one to each B), which converts M to a  doubly charged cation, while B atoms become singly charged anions. So, the MB2 formula can be expressed M21(B\\x00)2.31,32,40-45 The electron configurations vary depending on the donor prop erties of M, which produces a diversity of crystal structure types  and properties. The M-B bonds have ionic characteristics as a  result of  the donor-acceptor  interactions, but  they also have  covalent characteristics due to partial excitation of d electrons  and the formation of spd hybrid configurations. The tendency for B atoms to form sp2 and sp3 hybrids also affects properties.  However, hardness and brittleness are  lower  than the  corre sponding carbides because the B structural complexes combine sp3 hybridization with the lower-strength sp2 (and even lowerstrength s2p and sp) configurations, whereas the carbon atoms in carbides exhibit only sp3 hybridization.  As  shown in Fig. 2,  the crystal  structure of Group IV-VI  transition metal diborides  is primitive hexagonal  (AlB2-type, P6/mmm space group). The unit cell contains one MB2 formula unit. The structure is composed of layers of B atoms in 2D  graphite-like  rings or nets, which alternate with hexagonally  close-packed M layers. Each M atom is surrounded by six equi distant M neighbors in its plane and 12 equidistant B neighbors  (six above and six below the M layer). Each B is surrounded by  three B neighbors in its plane and by six M atoms (three above  and three below the B layer).  The unit-cell parameters and interatomic distances for dibor ides are  summarized in Table  II.  In general, B-B separation  controls the a-axis length. However, the a-axis length is also af fected by M-B contact. For ZrB2, which has the largest M atom (a/O3), which exof the diborides, the B-B distance is 1.83 A˚ the ‘‘normal’’ B-B distance (1.74 A˚ ) by 0.09 A˚ due to  ceeds  stretching of  the B-B bonds by Zr-Zr contact. Likewise,  the  smallest M atoms (Cr, V) lead to reductions in the B-B distance.  From crystal chemical considerations, the length of the a-axis is  a balance between two forces: (1) repulsion between atoms in the M layers and (2) attraction between atoms in the B nets.26,29,30  As a result, stable AlB2-type diborides do not form for M atoms smaller than Cr or larger than Zr. The B-B bond length for  minimum strain in the boron nets has been estimated to be B1.75A˚ , which is the value for TiB2.29 The M-B distance in diborides increases linearly with the M:B radius ratio, increasing from 2.30 A˚ for CrB2 to 2.54 A˚ for ZrB2. The M-B separation is equal to ða2 =3 þ c2 =4Þ1=2 , which is larger than the sum of the M and B radii (Table II). Generally,  the structural data indicate that bonding in the B nets controls  the lattice parameters in the AlB2-type structure. Because the B- B bonds are strong relative to the other bonds, increases in the a axis with increasing M size are minimal.  In contrast, no such  effect  is observed for the c-axis. Hence,  the c:a ratio increases  with increasing M atom size. Owing to separation of the close packed M planes by B nets,  the c-axis  is always  substantially  larger than 2RM. In summary, differences among diborides result from different M radii, which lead to variations in the in teratomic bond lengths. Larger changes are observed for the c axis than for the a-axis due to the relative bond strengths.  (2)  Structure-Property Relations  Hardness, bulk modulus, Debye temperature (YD), Tm, coefficient of thermal expansion (CTE), thermal conductivity (k), and formation ðDH o f Þ are some properties that are related to bond strength (cohesive energy). Generally, the combi enthalpy of  nation of bonds (M-M, B-B, and M-B) influences the material  properties. However, in some cases, a specific type of bond controls a property.31 For example, B-B and M-B bonds in dibor ides control hardness and thermal stability. Hardness therefore, a qualitative indicator of bond strength.47  is,  The thermal and elastic properties of diborides are summarized in Table III.48-54 The data indicate that the Group IV di borides have lower CTE and higher Young’s modulus and thermal conductivity than Group V diborides.26 The property  changes suggest that B-B bonds are strongest for Group IV at oms (Ti, Zr, and Hf) and the bonds weaken as atomic number  increases across a period of the Periodic Table. Grimvall and Guillermet46 found it useful to correlate cohesive energy and Tm trends to the average number of valence electrons per atom for  isostructural diborides. For example, the series ScB2, TiB2, VB2, CrB2, MnB2, and FeB2 have 9, 10, 11, 12, 13, and 14 valence electrons, respectively. With its 10 valence electrons (3.3 elec trons per atom), TiB2 possesses the highest melting temperature of the group. Higher or lower numbers of valence electrons re sult in lower melting temperatures.  Fig. 1.  A comparison of the melting temperatures of  the most refrac tory members of several classes of materials. Several borides, carbides, and nitrides have melting temperatures above 30001C and are considered  ultra high-temperature ceramics. For comparison, the melting temperature of Zr is B18501C and the melting temperature of Hf is B22271C.  Table I.  Summary of Some Structural, Physical, Transport,  and Thermodynamic Properties of ZrB2 and HfB2  Property  ZrB2  HfB2  Crystal system space group  prototype structure a (A˚ ) c (A˚ ) Density (g/cm3) Melting temperature (1C)  Hexagonal5  P6/mmm AlB2 3.17  Hexagonal6  P6/mmm AlB2 3.139  3.53 6.1195 32451 4899  3.473 11.2126 33801 48011  Young’s modulus (GPa)  Bulk modulus (GPa)  215 239 5.9 \\x02 10  212 287 6.3 \\x02 10  Hardness (GPa)  Coefficient of thermal \\x001) expansion (K Heat capacity at 251C (J \\x01 (mol \\x01 K)\\x001) Electrical conductivity (S/m)  \\x006 7  \\x006 7  48.28  49.57  1.0 \\x02 107 7 607  9.1 \\x02 106 7 1049  Thermal conductivity (W \\x01 (m \\x01 K) \\x001) Enthalpy of formation at 251C (kJ)  \\x00322.68  \\x00358.17  Free energy of formation at 251C (kJ)  \\x00318.28  \\x00332.27  1348  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  \\x0c', 'Hardness can also be related to electronic structure despite  wide differences  in published data. The hardness of diborides  decreases  as  the  atomic number  of  the metal  increases  for  Groups  III-VI  and Periods  4-6 of  the Periodic Table. The  Fermi  level  for ZrB2 completely occupied bonding states and free antibonding states.  is  located in a pseudogap between the  As a result, ZrB2 has the maximum stability and microhardness for its period. Likewise, TiB2 and HfB2, which are isoelectronic and isostructural, have the highest hardness among their peri ods. For Group V diborides, a considerable number of valence  electrons enter antibonding states, leading to decreases in bond strength and microhardness.35 Hardness HfB2 decrease as temperature increases.55 The bonding in diborides also influences  for TiB2, ZrB2, and  the anisotropy in  properties. Microhardness measurements for TiB2 conducted by Vehldiek52 did not reveal any significant anisotropy in hardness  along the aor c-axes. In comparison, the Young modulus showed a high degree of anisotropy in some diborides.53 Esti mation of stresses from strain-induced broadening of X-ray dif fraction (XRD) peaks revealed that the Young’s moduli of TiB2 and ZrB2 were reasonably isotropic, while NbB2 was extremely anisotropic, with lower values observed for the [00l] direction.56  This behavior  is observed in structures where  the bonds are  weaker in one direction than for other directions. Based on line  broadening studies, it was concluded that TiB2 had strong bonding in all directions, bonding in ZrB2 was isotropic, but weaker than in TiB2, and that NbB2 was highly anisotropic with weak bonding between (00l) planes. Post et al.30 were first to show that diborides with metals of  large radii had the highest melting temperatures. Because the B-  B separations are the highest and the B-B bonds are the weakest  for the highest melting temperature diborides, the authors came  to the conclusion that  the M-B bond strength was responsible  for melting temperatures. The CTEs of various diborides have been measured.57-59 Using XRD analysis up to 16001C, Keihn and Keplin57  showed  that ZrB2 and HfB2 had lower CTE anisotropy than the diborides of Ti, Nb, and Ta. The authors concluded that polycrys talline ZrB2 and HfB2 ceramics would develop lower stresses during rapid cooling than other diborides. With the ex internal  ception of CrB2, the lattice parameters and CTEs of the diborides had similar temperature dependences for the aand c-axes,  which indicates little anisotropy in the bond strength in the two  directions. The CTE in the c-direction decreases with increasing  M radius, which can be correlated with an increase in the M-B  bond strength with increasing M size. The CTE in the a-direc tion does not  change  significantly with increasing M radius.  Taken together, the results can be interpreted to indicate that the  B-B bonds determine  the  cohesive  forces  in the a-direction,  while weaker M-B bonds control cohesive behavior direction.59  in the c Fig. 2.  Projections of  the AlB2-type structure. Pictures are taken from the Crystal Lattice Structures Web page, provided by the Center for Computational Materials Science of the United States Naval Research Laboratory.  Table II.  Unit Cell Parameters and Interatomic Distances in Diborides26,30  Metal  boride  Size or length (A˚ )  c/a  c  a  R0  M  B-B  M-B  M þ RB 2R0  ZrB2 HfB2 TaB2 NbB2 TiB2 CrB2  1.114  3.530  3.169  1.61  1.830  2.54  2.48  1.105  3.470  3.141  1.58  1.813  2.51  2.46  1.041  3.225  3.097  1.49  1.788  2.41  2.36  1.05  3.27  3.11  1.48  1.79  2.43  2.36  1.064  3.228  3.028  1.47  1.748  2.38  2.34  1.030  3.066  2.969  1.30  1.714  2.30  2.18  R0 M is the metal radius and RB is the radius of boron.  May 2007  Diborides of Zirconium and Hafnium  1349  \\x0c', '(3)  Solid Solubility  Mutual continuous solid solutions form among Group IV and V diborides. Because their radii differ by o10%, HfB2 and ZrB2 form continuous solid solutions with TiB2, NbB2, and TaB2. A larger radius difference leads to limited solubility between these diborides and CrB2.25,30 The activation energy for the formation of diboride solid solutions is essentially the same as the activa tion energy for M diffusion. For example, the activation energies  for the formation of solid solutions in the ZrB2-TiB2 and TiB2- NbB2 systems are 175 and 112 kJ/mol, respectively, which correspond to the activation energies for metal diffusion in these  systems. In comparison, the NbB2-CrB2 system exhibits higher values for activation energy (400 kJ/mol), which can be attrib uted to the large disparity in atomic radii between Nb and Cr  (13%) compared with Zr and Ti (9%) or Ti and Nb (1%). Based  on activation energy data, solid solution formation appears to  depend on diffusion of M atoms, not  the strongly bonded B  atoms.  Some diboride solid solutions show deviations from ‘‘rule of  mixture’’ predictions for properties such as hardness and elec trical conductivity, although some changes are only observed  over limited composition ranges that are specific to each system.  For example, TiB2-ZrB2 solid solutions exhibited a maximum in hardness between 30 and 70 mol% ZrB2. The hardness at 50 mol% ZrB2 was 24 GPa compared with values of 18 GPa for TiB2 and 12 GPa for ZrB2 as reported in the study.60 For TiB2- ZrB2 ceramics synthesized by a combustion synthesis-microwave process followed by hot pressing (HP), the highest values  of relative density (99%), flexure strength (680 MPa), (7.3 MPa \\x01 m1/2), and hardness toughness served at 20% ZrB2.61 In another study, TiB2-ZrB2 ceramics containing 40 mol% ZrB2 had the highest hardness (30.4 GPa), fracture toughness (3.87 MPa \\x01 m1/ flexure strength (354 MPa), 2), Young’s modulus (449 GPa), and fracture energy (16.5 J/ m2).62 Local maxima in hardness have also been observed in systems such as TiB2-NbB2 (at B38 mol% NbB2),63 NbB2-CrB2 (30 GPa at 80 mol% CrB2),64,65 and TiB2-CrB2 (at B20 mol%CrB2).66 The lattice constants of diboride solutions show a linear de fracture  (27 GPa) were ob other  pendence on composition in many systems. However, based on  bonding considerations discussed above, c-axis parameters are  more strongly affected by changes  in composition than a-axis  parameters. For TiB2-WB2 solid solutions with WB2 contents  between 40 and 60 mol%, the a parameter decreased only slight ly as WB2 content increased, whereas the c parameter exhibited a significant decrease.67 For TiB2-CrB2 and TiB2-ZrB2 solutions, Zdaniewski68 found that the degrees of lattice contraction  (for Cr) or expansion (for Zr) were smaller than expected based  on the sizes and concentrations of  the M atoms. This suggests  stronger interactions between the M and B atoms than would be  expected from the M size alone and is likely due to differences in electron configurations and bond strength.35,40 Changes  in  lattice anisotropy and elastic  strain energy in solid solutions  may also affect grain-boundary cohesion in polycrystalline ceramics.68  If  true,  then knowledge of  lattice parameter changes  with composition could be used to improve  the mechanical  behavior of diboride ceramics by using compositional changes  to induce beneficial stresses (for example, surface compression  stresses). Fendler et al.69 studied thermal expansion of TiB2-CrB2 and TiB2-WB2 solid solutions using high-temperature XRD analysis. The CTEs of TiB2-CrB2 did not follow predictions from rule of mixtures calculations. The c-parameter exhibited an unex pected increase with temperature. For TiB2-WB2 tions, the CTE was lower than TiB2 at low temperatures. Using experimental hardness values obtained from solid solutions as a function of composition,61,64,66 NSWCCD research solid solu ers calculated the B-B bond lengths for compositions showing  the maximum hardness in TiB2-CrB2, NbB2-CrB2, and TiB2- ZrB2 solid solutions. The B-B bond lengths were calculated to be 1.74, 1.73, and 1.75 A˚ , respectively. The values were similar to 1.75 A˚ , which Hughes and Hoard29 suggested as the B-B bond  length for minimum stress in the boron net. The result  is also  consistent with hardness values from pure diborides, which were  higher for borides with the largest M atoms.  III.  Synthesis  The diborides can be synthesized by a variety of  routes. This  section focuses on three main synthesis  routes:  (1)  reduction  processes, (2) chemical routes, and (3) reactive processes.  (1)  Reduction Processes  Several  reduction processes are used to synthesize diborides.  Carbon and boron are the most common reducing agents, but  boron carbide (B4C) or aluminum (Al) can also be used as well as combinations of reducing agents.70 Examples of the reactions  used to synthesize diborides are shown in Table IV (Reactions  (1)-(5)). These reactions were written for ZrB2, but analogous processes produce HfB2 or other diborides. Carbothermal reduction is used to produce ZrB2 and HfB2 commercially.71-73 The process is strongly endothermic; Reaction (1) has an enthalpy of reaction ðDH o rxn Þ of 1475.6 kJ at 298 K. As a result, Reaction (1) is only thermodynamically favorable rxn < 0Þ above B15001C (Table IV). Based on reaction (i.e., DGo favorability and the heat absorbed, temperatures of B20001C are generally employed to synthesize diborides.71,74 The reac tions also produce significant volumes of various gases, which  must be removed for  the reactions  to proceed to completion.  Excess B2O3 is often added to promote formation of diborides over carbides. Along with B2O3 and carbides, carbon is also a common impurity in the final powder.70  Table III.  Thermal and Elastic Parameters of Transition Metal Borides26  Boride  CTE  (ppm/K)  YD (K)  E  (GPa)  HV  (GPa)  k W \\x01 (m \\x01 K)  \\x001  Calc.  from a  Calc  from TM  TiB2 ZrB2 HfB2 VB2 NbB2 TaB2 CrB2  4.8  1100  970  530  13  20.6  6.2  765  730  420  15  18.9  6.6  550  580  —  —  16.6  8.0  880  880  340  —  13.1  8.2  701  720  —  —  6.9  8.5  545  570  280  —  6.0  10.5  726  780  220  13  31.8  Table IV.  Examples of Reduction Reactions that can be Used to Synthesize Borides and the Change in Free Energy of Reaction rxn Þ as a Function of Temperature  ðDG o  Reactions  Category  Example  ðDGo  rxn Þ (kJ)  (1)  Carbothermal  ZrO2 ðcÞ þ B2O3 ðl Þ þ 5CðgÞ ! ZrB2 ðcrÞ þ 5COðgÞ ZrO2 ðcÞ þ 4BðcÞ ! ZrB2 ðcÞ þ B2O2 ðgÞ 3ZrO2 ðcÞ þ 3B2O3 ðl Þ þ 10Alðl Þ ! 3ZrB2 ðcÞ þ 5Al2O3 ðcÞ 7ZrO2 ðcÞ þ 5B4CðcÞ ! 7ZrB2 ðcÞ þ 3B2O3 ðgÞ þ 5COðgÞ 2ZrO2 ðcÞ þ B4CðcÞ þ 3CðgÞ ! 2ZrB2 ðcÞ þ 4COðgÞ  1431-0.803T  (2)  Borothermal  301-0.178T \\x00241210.495T 1378-0.924T  (3)  Aluminothermal  (4)  Boron Carbide  (5)  Combined  1134-0.668T  1350  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  \\x0c', 'Reduction with B4C have also been used to produce diborides and diboride composites.70-73,75 Like other reduction processes,  the reactions between oxides and B4C are endothermic. However, these reactions become favorable at temperatures lower  than the  corresponding reactions with carbon. For example, Reaction 4 becomes favorable around 12001C whereas Reac(2), and (5) are only favorable above 14001C (Table  tions (1),  IV). Processes similar to Reaction (4) offer the additional benefit  of reduced levels of carbon, B2O3, and oxide impurities in the reaction products. Reaction (4) has also been exploited to pro mote densification of B4C and ZrB2 based on the ability of B4C to react with oxides and thereby reduce the amount of oxide impurities in non-oxide ceramics.76,77  (2)  Chemical Routes  Chemical  routes used to produce diborides  include solutions,  reactions with  boron-containing  polymers,  and  pre-ceramic  polymers. Nanocrystalline ZrB2 and HfB2 have been synthesized by reacting anhydrous chlorides with sodium borohydride (NaBH4) above 5001C under pressure.78,79 The overall process is described by Reaction (6), but the reaction may proceed by a  vapor phase mechanism that  involves  the decomposition of  NaBH4 to borane (BH3) and its subsequent reaction with a gaseous chloride. Diborides prepared by these routes can have crystallite sizes as small as 10-20 nm78,79:  HfCl4 ðgÞ þ NaBH4 ðgÞ ! HfB2 ðcÞ þ 2NaClðcÞ þ 2HClðgÞ þ 3H2 ðgÞ  (6)  Titanium diboride has also been prepared by thermal decom position of  titanium borohydride (Reaction (7)). Powders pre pared by this method were also nano-sized with diameters of 100-200 nm80:  2TiðBH4 Þ3 ðcÞ ! 2TiB2 ðcÞ þ B2H6 ðgÞ þ 9H2 ðgÞ  (7)  Nano-crystalline oxides produced by precipitation could be  converted to diborides using reduction processes (Table IV). Sacks et al.81 synthesized nano-crystalline ZrC and HfC by car bothermal reduction of ZrO2 and HfO2 produced by a sol-gel process. Similar processes should also work for nano-crystalline  diborides, but, to date, only conventional micro-sized oxides have been converted to diborides.82 A pre-ceramic polymer has been directly converted to a refractory diboride,83 but  the pol ymer was used as a binder for commercial ZrB2 powder, not to produce a stand-alone coating or a monolithic ceramic. Finally,  diboride-based composites have been prepared by combining  conventional diboride powders with SiC-bearing pre-ceramic polymers such as polycarbosilane.84,85 From these reports, a va riety of chemical synthesis methods are available, but the proc essing of these powders into coatings or ceramics has not been  fully explored.  (3)  Reactive Processes  The simplest reaction synthesis method for producing diborides  is the reaction of elemental precursor powders:  Zr þ 2B ! ZrB2  (8)  Hf þ 2B ! HfB2  (9)  The use of direct  reactions dates back over 100 years, al though pure diborides were not produced until nearly 50 years later due to the difficulty in obtaining high-purity boron.70 The  high oxygen affinity of both Zr and Hf means that the reactions  must be carried out in inert or reducing atmospheres to prevent  the formation of oxide impurities.  Reactions (8) and (9) are both extremely exothermic rxn ¼ \\x00323 and \\x00 328kJ, respectively, at 298 K) and favor(DH o  able at all temperatures.10,86 If allowed to proceed in an uncon trolled manner,  the heat can ignite a self-propagating reaction  and generate temperatures high enough to promote local melting of the Zr (Tm 5 B18501C) or Hf (Tm 5 B22251C).87 Thus, Reactions (8) and (9) can be used to synthesize ZrB2 and HfB2 by self-propagating high-temperature synthesis (SHS), which is also known as combustion synthesis.88-90 The high heating and  cooling rates associated with SHS produce high defect concen trations  in the resulting diborides, which have been linked to powders.91 The  improved  sinterability  of  SHS-derived  SHS  method also lends itself  to preparation of diboride-containing  composites or diborides doped with additives such as sintering aids.92,93 In addition, a great deal of work has been done on SHS  of TiB2 ceramics, all of which is directly applicable to ZrB2 and HfB2. In contrast to SHS, recent research at UMR has employed slow heating (B11C/min) and extended isothermal holds (6 h at 6001C) to react fine powders of Zr and B without the ignition of SHS reactions.94 Under controlled heating, Zr and B react to low as 6001C. The  form phase-pure ZrB2 at ZrB2 crystallites produced by controlled reactions were the same size as the Zr precursor particles. Hence, reduction of the start temperatures as  ing Zr particle size before reaction reduced the size of  the re sulting ZrB2. Analysis showed that the crystallite size of ZrB2 formed from Zr attrition milled for 240 min was B10 nm. Par ticles coarsened significantly as temperature increased and grew to more than 1 mm at 16501C (Table V). Coarsening in non oxides ceramics, including diborides, is promoted by oxygen present as oxide impurities on the particle surfaces.3,77,95,96 A  significant amount of oxygen (2.4 wt%) was detected in ZrB2 produced by Reaction (8), which was higher than reported for  conventional powders (0.9% for as received H. C. Starck Grade  B powder, 2.0 wt% after attrition milling for 2 h in hexane).  IV.  Densification  The need for dense diboride ceramics was  initially driven by  nuclear applications that required high neutron absorption cross sections. Later, research was driven by aerospace applications.71  Without  sintering additives,  full density has, historically, only  been achieved by HP. From the early 1950s to the 1970s, con siderable  effort was directed at densification of commercially  available, non-reactive diboride powders. Research focused on  ways to activate sintering so that larger components with more  complex shapes could be produced as HP has limits related to  both size and geometry. Metallic additives and liquid-phase sin tering aids were used to enhance densification. Recent  studies  have  again focused on additives  that promote densification,  along with some modern densification techniques. As a result,  progress has been made in understanding the sintering of di borides while minimizing second phases that may be detrimental  to their high-temperature properties. Reports of fully dense diborides, produced by reactive HP date back to 1951,70 indicating  the viability of this approach as well.  This section discusses four densification methods: (1) HP; (2)  pressureless sintering; (3) reactive routes; and (4) spark plasma  Table V.  Crystallite/Grain Size and Density for ZrB2 and ZrB2-SiC Produced by Reactive Hot Pressing of Zr-B and Zr-B-SiC Powder Mixtures  Processing temperature (1C)  ZrB2  ZrB2-SiC  Crystallite  size (nm)  Relative  density (%)  Crystallite  size (nm)  Relative  density (%)  600  10  36  10  37  1000  50  35  50  37  1450  600  34  300  40  1650  1000  40  500  95  May 2007  Diborides of Zirconium and Hafnium  1351  \\x0c', 'sintering. Induction zone melting has also been used to densify ZrB2 and HfB2.55,97 However, this process was not (B0.25-0.6 because the specimen size cm) only allowed for (grain size of B300 mm) and  reviewed  limited microstructural analysis  property measurements (microhardness).  (1)  HP  Because of strong covalent bonding and low self-diffusion, high  temperatures and external pressures are required to densify diborides.2 Results of early HP studies on commercially available ZrB2 and HfB2 powders are summarized in Table VI,7,12,98-104 which includes details of the starting materials, HP conditions,  and final densities. Most of the diborides were presumed to be al.99  stoichiometric  (B:M 5 2.0). However, Kalish et  studied  boron-deficient  compositions  (e.g., ZrB1.98 composition was nominally in the  and Hf1.82). The single-phase ZrB2 thus, should not have affected the HP conditions or  ZrB1.98  field and,  the resulting microstructure. Metal additions used to produce  HfB1.82  should place  the  composition in the HfB2-HfB twophase field. It is not clear how this change in composition may  have affected densification and the resulting microstructure, but the HP temperature (18001C) was well below the solidus temthe HfB2-HfB binary (B21001C). Because no liquid should form, densification would still have to proceed by a  perature for  diffusion-controlled (as opposed to a liquid phase) mechanism  and would be expected to occur at a temperature similar to the  stoichiometric diboride. HP of both ZrB1.98 18001C required the application of a high  and HfB1.82 at (B827 pressure  MPa) to reach full density.  In historic studies, nominally stoichiometric ZrB2 and HfB2, without additives, have only been densified by HP at 20001C or higher with pressures of 20-30 MPa,105 or at reduced temperatures (1790-18401C) with much higher pressures (800-1500 MPa).99,102 Recent studies (also included in Table VI) have pro duced similar results. Nearly phase pure ZrB2 and HfB2, with average starting particle sizes of 5-10 mm, require HP above 20001C to achieve full density.104 For example, HfB2 with a nominal starting particle size of 10 mm was o95% dense after HP at 21601C and 27.3 MPa for 180 min.12  Recent research at UMR showed that commercial ZrB2 with a B2 mm starting particle size (Grade B, H. C. Starck, Newton, MA, 499% purity with 0.89% O) can be further comminuted  by standard techniques (e.g., attrition milling) to enhance den sification. Particle size reduction reduced the HP temperature necessary to achieve full density to 19001C (pressed for 45 min at 32 MPa).7 Analysis revealed no porosity, but B1.9 vol% WC  was present. The WC was incorporated during attrition milling  due to wear of the WC-based milling media. The microstructure  contained both equiaxed and slightly elongated ZrB2 grains with an average grain size of B6 mm. Minimization of grain growth  was attributed to lower HP temperatures and reduced starting particle sizes. This result can be contrasted with 10-20 mm av erage grain sizes that are common in diborides processed from  coarser ZrB2  and HfB2  powders with  higher HP tempera tures.12,100,101,106  Research in the late 1960s on metal (Zr, Hf, Cr, Y, and Al)  and ceramic (SiC, MoSi2, and TaB2) additions to ZrB2 and HfB2 revealed the superior densification, and more importantly the  improved oxidation and thermal stress resistance (TSR), of diborides containing SiC, C, or SiC plus C.107,108 Fenter100  re viewed many compositions,  the most  important of which were  ZrB2 and HfB2 with 20 vol% SiC or 18 vol% SiC and 10 vol% C. Subsequent to these key studies, recent research using com mercial ZrB2 and HfB2 powders has typically included non-reactive additives (e.g., SiC) or liquid phase forming additions  (e.g., Ni, Si3N4, or MoSi2) to improve densification. The high HP temperatures and pressures of historic studies99,102 are not  necessary with the finer  starting powders and additives  that  minimize grain growth, produce a liquid phase, or result in solid  solution formation. In recent studies, both ZrB2 and HfB2 have been hot pressed to near full density by adding Ni,109,110 Ta5Si3,119  Si3N4,110,111 AlN,110,112 MoSi2,113-115  SiC,7,116-118 additives.109,120-123  and  other UHTCs,  or  combinations  of  Many of these additives reduced the densification temperature of ZrB2, with Ni reducing it to as low as 16001C.120 However, even with these additions, pressures of 20-50 MPa were re quired.  (2)  Pressureless Sintering (PS)  PS of diborides would enable the fabrication of components to  near-net  shape  using  standard  powder  processing methods.  Compared with HP,  sintering  is more  efficient  and reduces  costs, potentially opening the door for other applications. The  two approaches typically used to enhance densification are re duction of  starting particle size and the use of  sintering aids.  Potential sintering aids can be classified into categories including  liquid phase formers, solid solution formers, and reactive agents.  Some discrepancy exists for the onset temperature for densification of diborides. Temperatures as low as 13001-15001C were reported in historic studies,98,102 while more realistic temperatures of B17501C have been reported recently.106 The dis crepancy is likely due to the purity of the powders. Densification  of pure diborides is due to grain boundary and volume diffusion, which are not appreciable until 18001C or higher.2 The  presence of oxide or metallic  impurities  increases  the driving  force for densification, but also enhances grain and pore growth,  Table VI.  Hot Pressing Conditions and Densities of Hot-Pressed ZrB2 and HfB2  Boride  Particle size (mm)  Hot pressing conditions  Final density (%)  References  Temperature (1C)  Pressure (MPa)  Time (min)  Historic studies  ZrB2  7  2100  B28 B827 B17  80  B95 r100 99.2 B9572 r100 r100 r100 98.0  Chown99 Kalish et al.100 Fenter101 Andrievskii et al.102 Kalish and Clougherty103 Kalish and Clougherty103 Kalish et al.100 Fenter101  ZrB1.89  —  1800  —  ZrB2 ZrB2 HfB2 HfB2  —  2000  260  —  2200-2300  29-39 B793 B1,551 B827 B23  20-30  5  1840  10  5  1790  10  HfB1.82  5  1800  —  HfB2 Recent studies  —  2200  200  ZrB2 ZrB2 HfB2  6  1900  30  30  86.5  Bellosi and colleagues104,105 Cutler9 Opeka et al.12  o2  1900  32 B27  45  99.8  B10\\x03  2160  180  90-95  [—],  information not disclosed in the reference. \\x03\\x00325 mesh (nominally B10 mm on average) HfB2 powder from Cerac Inc., Milwaukee, WI.  1352  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  \\x0c', 'May 2007  Diborides of Zirconium and Hafnium  1353  which limit the maximum achievable density to 95% or less. In  particular, B2O3 enhances grain growth and inhibits densification by promoting coarsening at temperatures below 18001C.124,125  Historic studies implied that phase-pure diborides could not  be pressurelessly sintered to full density since  the hexagonal  crystal structure allowed for anisotropic grain growth and en trapped porosity (i.e., coarsening was more favorable than densification). However, Baumgartner and Steiger126 successfully sintered TiB2 to 499% density at 20001-21001C using submicrometer starting powders. While fine particle size enhanced densification, Baik and Becher3 later showed that the other key  to pressureless densification of TiB2 was the total oxygen conto r0.5 wt% minimized tent. Reducing the oxygen content evaporation-condensation, which allowed densification to pro ceed without significant grain or pore growth.  Historic studies on PS of ZrB2 and HfB2 are limited, and many were carried out nearly 40 years ago.127 However, in creased densification rates were demonstrated through additions of transition and refractory metals.127 Meerson and Gorbunov128 used 0.7 wt% C, along with ZrO2 and B, nearly fully dense ZrB2. They hypothesized that the C reduced ZrO2, which subsequently reacted with the B to produce active ZrB2 nuclei on the surface of larger ZrB2 grains. Cech et al.129 added metals to accelerate densification. They reported B100% density for ZrB2 sintered at 22001C for 1 h in argon when adding 1 wt% Re, or 0.5 wt% Cr 10.5 wt% Ti. Lattice parameter  to produce  measurements after  sintering showed metal atom substitution  onto Zr sites in the ZrB2 lattice. They hypothesized that the local contraction of the lattice affected the surface and volume free  energies, effectively increasing the driving force Coble and Hobbs124 showed similar et al.130 used W additions  for TiB2. Kisliy to ZrB2. They reported significant densification, along with evidence for substitution of some W  for  sintering.  results  into the ZrB2 lattice. Gropyanov131 also noted that  intensive mechanical milling,  vibratory milling in this case, enhanced densification of ZrB2. He hypothesized that the milling not only increased the surface  area, but also increased the number of point defects near  the  surface of the particles, enhancing the driving force for densifi cation by grain boundary diffusion.  In a recent study of pressureless sintering, Kida and Segawa132,133 sintered ZrB2 to 95% relative density without external pressure. However, densification was only possible with signif icant amounts of BN (5 wt%), AlN (15 wt%), and SiC (5 wt%). Quabdesselam and Munir134 found that fine TiB2 produced by SHS did not have improved densification compared with pow ders produced by carbothermal reduction. Their results demon strated that powder processing alone does not activate densification. Mishra et al.88 sintered ZrB2 with a nominal particle size of 2.8 mm, produced by SHS, to a density of 94% at 18001C for 1 h in flowing argon. They noted a small amount of  ZrO2 during microstructural analysis, but and Cr impurities promoted local melting, which improved densification. In a subsequent study, Mishra et al.93 focused on Fe  concluded that Fe  and Cr additions to ZrB2. Although Cr resulted in swelling and cracking, the addition of up to 10 wt% Fe improved density to B93% at 18001C. While transition metal additives reduce the  densification temperatures for diborides, they also cause marked  reductions in melting temperature, hardness, and high-temperature strength.109,128,129  Modern ZrB2 can be sintered to near metal additions. Sciti et al.135 sintered ZrB2 (Grade B, H. C. Starck, Germany) to full density at 18501C by adding 20 vol%  full density without  MoSi2, which led to liquid phase formation. Research at UMR has shown that the same ZrB2 can be densified using reactive additives.77,136 With small amounts of WC (B2 vol%), ZrB2 was sintered to 498% density at 21501C for 540 min. Analysis  by XRD indicated that W and C were incorporated into the  ZrB2 lattice after reaction. Similar to the work of Baik and Becher3 in TiB2, densification of ZrB2 was activated when the oxide impurities (B2O3 and ZrO2) were removed from particle  Fig. 3.  Relative density as a function of sintering temperature for ZrB2 (as-received and attrition-milled powders), with and without 4 wt% B4C additions. The addition of B4C led to higher relative densities at all sintering temperatures for both attrition-milled powder, achieving 100% of  the theoretical density for milled ZrB2 with 4 wt% B4C sintered at 18501C.  surfaces, which minimized coarsening. Thermodynamic predic tions followed by XRD analysis confirmed that ZrO2 could be removed with either WC or B4C. However, B4C reacted with lower temperatures (B12001C) than WC (B18501C), ZrO2 at which reduced the densification temperature (Fig. 3). The mi crostructures of ZrB2  sintered with B4C (Fig. 4) confirm that  Fig. 4. Microstructure of ZrB2 sintered at 18501C for 1 h in vacuum. Specimens were prepared from attrition-milled powder with (a) 2 wt%  B4C and (b) 4 wt% B4C additions.  \\x0c', 'dense ceramics 18501C.  can  be prepared  at  temperatures  as  low as  (3)  Reactive Routes to Densification  Synthesis and densification can be combined into single-step,  in situ reaction and densification processes. Either elemental re actions (Reactions (8) and (9)) or reduction processes (Reactions  (1)-(5)) can be used to form diborides by reactive HP (RHP). An extremely fine crystallite size (B10 nm) was achieved for ZrB2 by reacting Zr and B that were attrition milled.94 The fine crys tallite  size  should enhance  the driving force  for  sintering, as  densification is driven by minimization of surface free energy. However, HP at 21001C was required to achieve a relative den sity of 99% for ZrB2, despite its extremely fine crystallite size, possibly due to grain coarsening below the final densification temperature.94 For comparison, ZrB2 ceramics produced from commercially available powders with submicrometer-sized particles could be hot pressed to 499% density at 19001C.137 Anal ysis of scanning electron microscopy (SEM) images revealed an average grain size of B12 mm for ZrB2 produced by RHP at 21001C, compared with B6 mm for ZrB2 produced by conventional HP at 19001C.137 The crystallite size and density data in  Table V confirm significant coarsening in the RHP material. As has been reported for other non-oxide ceramics such as TiB2,3 B4C,138 and SiC,139 oxygen impurities in ZrB2 inhibited densification by promoting coarsening at temperatures below which  densification could occur.  Combined in situ synthesis and densification processes have also been developed for ZrB2-SiC94,140 and HfB2-SiC.141 These processes employ reactions of elemental precursors (Zr, B, SiC)142  and  as well  as more  complex  displacement  reac tions140,141,143,144:  2Zr þ Si þ B4C ! 2ZrB2 þ SiC  (10)  2 þ x ÞHf þ 1 \\x00 x ÞSi þ B4C ! 2HfB2 þ 1 \\x00 x ÞSiC þ xHfC  ð  ð  ð  (11)  The presence of SiC had a dramatic effect on both the den sification temperature and the resulting microstructure. Using Zr, B, and SiC as precursors, a relative density of B95% was achieved by RHP at 16501C (Table V). Not only was the densification temperature (16501C) reduced significantly compared with ZrB2 produced by RHP (21001C) or conventional HP (19001C), but the ZrB2 grain size was also much smaller. Analysis of the microstructure (Fig. 5) revealed an average grain size of B0.5 mm for ZrB2-SiC produced by RHP compared with B12 mm for dense RHP ZrB2. The reduction in processing temperature for ZrB2-SiC by RHP has two likely causes: (1) mini mization  of  oxide  impurities  (B2O3,  ZrO2)  and  (2)  fine  crystallite/particle sizes.  (4)  Spark Plasma Sintering (SPS)  SPS has recently been applied to densify UHTCs114,115,122 and other ceramics.145-147 The SPS process consists of simultaneous ly applying a uniaxial load and a direct or pulsed electric current  to a powder compact. SPS is similar to HP. However, in place of  indirect heating, the applied electrical field heats the die and the  powder compact (if the powder is electrically conductive). Rapid heating rates (hundreds of 1C/min) facilitate rapid densification,  which minimizes grain growth. As an example, Monteverde et al.148 used SPS to achieve full density for HfB2130 vol% SiC by heating at 1001C/min to 21001C with a 2 min hold and at 30 MPa. The entire SPS cycle took B30 min,  including heating  and cooling, which resulted in a uniform microstructure with a mean grain size of B2 mm. Medri et al.122 compared ZrB2 containing 30 vol% ZrC and 10 vol% SiC prepared by HP and SPS. Despite the addition of  3.7 vol% Si3N4 as a sintering aid, the HP specimen had a maximum densification rate of only 1.3 \\x02 10\\x005 s\\x001 at 18701C, ulreaching only B90% density. timately specimen had a nearly constant densification rate of 2 \\x02 10 In contrast, the SPS \\x004 s over a broad temperature range (17501-20501C), resulting in near full density at 21001C in o60 min. Similar results were also observed in HP versus SPS comparison studies for HfB2115 vol% MoSi2,114 ZrB2115 vol% MoSi2,115 and HfB2130 vol% SiC12 vol% TaSi2.149  \\x001  V.  Microstructure and Mechanical Properties  Besides melting temperature, the properties that make ZrB2 and HfB2 attractive for high-temperature structural applications are strength at elevated temperatures and TSR. This section focuses  on ZrB2 and HfB2, as well as diborides with SiC or MoSi2 additions, from recent studies. The subsections that follow discuss  the strength, modulus, hardness, toughness, and TSR of dibor ides from recent studies.  Results from historic studies are not discussed in this section  because the large starting particle size of the powders, higher impurity levels (40.25 wt% C or Fe), or densification temperatures (420001C), which resulted in exaggerated grain growth.2  Large grains reduced mechanical strength, due,  in part,  to re sidual thermal stresses that result from thermal expansion anisotropy.150 For anisotropic materials, a critical grain size exists avoided.151 While values  below which microcracking  can be  have not been established for ZrB2 or HfB2, the values should be similar to the critical grain size for TiB2, which is B15 mm.152 Hence, microcracking may influence the observed properties of large grained (i.e., 415 mm) diborides. In addition, many studies  used additives such as transition metals that reduce strength at elevated temperature (s1000 o 0.3sRT).103,104,109,110  (1)  Strength of Diborides  Like most ceramics, higher strengths are reported for diborides  with  finer  grain  sizes.7,114,121,153 The  dependence  of  flexure  strength on grain size (Fig. 6) demonstrates  that  strength in creases as grain size decreases for ZrB2 and HfB2 based ceramics, with and without MoSi2 or 135,140,141,148,154 As indicated in the legend, the plot includes di SiC.7,11,99-101,113-116,118,121,  borides produced by HP, RHP, SPS, and PS. Most strengths  were measured in four-point bending, but  some authors used  three-point bending. Further, specimen size varied, mostly be tween U.S. and European standards, and, thus, the strengths are  not directly comparable. However, strengths,  in general, show \\x001/2) as would an inverse square root relation with grain size (GS be expected for ceramics free from other, larger flaws. A line in Fig. 7 is intended to highlight an approximate GS\\x001/2 relationship, not as a fit to the data.  Fig. 5.  A polished,  thermally  etched  cross  section  of  a ZrB2-SiC ceramic prepared by the reactive hot pressing of Zr, B, and SiC at 16501C. SiC is the darker phase.  1354  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  \\x0c', 'May 2007  Diborides of Zirconium and Hafnium  1355  Fig. 6.  Room temperature flexure strength as a function of grain size  from the diboride literature.  Fig. 8.  Room temperature flexure strength as a function of the volume  percent additive (e.g., SiC, MoSi2, plus others). The shaded area is intended to indicate that all levels of additives increase the mechanical  strength, but that roughly 10 vol% may be enough to maximize strength.  Fabrication of diborides with fine grain size and high strength requires processing below 20001C (e.g., HP, RHP) or  limited  time at higher temperatures (e.g., SPS). Second phases such as  SiC, MoSi2, ZrC, or HfN that improve densification may also inhibit grain growth and improve strength.2,100,113,121 Research  at UMR has shown that hot-pressed ZrB2 (no additive) has a strength of 565 MPa (Fig. 7).7 The addition of 10, 20, or 30 vol% SiC or MoSi2 reduced the average grain size to B2-3 mm, which increased strength to 700-1000 MPa.7,113 Considering that the ZrB2110 vol% SiC composition was only B93.2% dense while the other compositions were fully dense,7 the results  indicate that 10 vol% of a second phase may be enough to limit  grain  growth  and maximize  strength  at  room temperature.  Room temperature  strengths  for diborides produced by HP,  SPS, and RHP support this conclusion (Fig. 8), particularly the  work of Monteverde.100,103,114,115,117,118,121,123,124,141,148,149,155  Although ZrB2 grain size is a factor, strength of fine-grained ZrB2-SiC is not controlled by the size of the ZrB2 grains.153,156 Rezaie et al.153 used linear elastic fracture mechanics (Griffith criteria) to show that the critical flaw size  the room temperature  correlates strongly to the SiC particle size in ZrB2-30 vol% SiC mismatch between SiC (aavg. B4 \\x02 10\\x006 K\\x001) and HfB2 ceramics. This conclusion is consistent with analysis of the CTE B8 \\x02 10\\x006 K\\x001) that results in GPa-level residual stresses that may produce microcracks at HfB2-SiC interfaces.141 The CTEs  (aavg.  of ZrB2 single crystals were reported to be 6.9 \\x02 10 \\x006 K \\x001 along the c-axis and 6.7 \\x02 10\\x006 K\\x001 along the a-axis.150 The difference in CTE values between the two axes within the hexagonal crystal structure results in a B95 MPa thermal stress within polycrys talline ZrB2 when cooled from a typical processing temperature (19001C).157 The addition of a-SiC (6H) particles results in  larger thermal stresses in ZrB2-SiC than in polycrystalline ZrB2 due to the lower CTE of SiC (4.7 \\x02 10 \\x006 and 4.3 \\x02 10 K\\x001 along the cand a-axis, respectively).158 An Eshelby analysis (Eq. (12)) of ZrB2 containing SiC-particulates, assuming spherical SiC inclusions, predicts  compressive  stresses  \\x006  radial  within the ZrB2 matrix to be as high as 2.1 GPa and tangential tensile stresses of 4.2 GPa at ZrB2-SiC interfaces159:  sradial ¼ \\x002 \\x02 stan  ¼  ðam \\x00 ai ÞDT ð1 \\x00 2ni Þ Ei þ ð1 \\x00 nm Þð2Em Þ  \\x12  \\x13  R r þ R  3  (12)  where m denotes the matrix, i denotes the inclusion, a is the CTE, n is Poisson’s ratio and E is Young’s modulus. The stresses  decrease as distance  from the  inclusion (r)  increases, and in crease as the inclusion size (R) increases.  Based on this analysis, smaller SiC particles reduce the ten dency for microcracking and the nucleation of larger, critical flaws.151 This conclusion is also supported by the strong corre lation between the room temperature flexure strength of ZrB2- 30 vol% SiC ceramics and the average SiC grain size, whereas  strength does not correlate to average ZrB2 grain size (Fig. 9). The data for Fig. 9 were obtained from specimens hot pressed at  various temperatures and times using a single SiC starting particle size,153 and from specimens hot pressed using SiC with different particles sizes at the same HP temperature.156 Analysis of  the critical flaw size along with maximum ZrB2 and SiC grain sizes, confirmed that the maximum SiC grain size in ZrB2-SiC ceramics is the strength limiting factor.156  Second phase additions reduce densification temperature and  grain size, but the resulting ceramics may experience mechanical elevated temperature.117 Modern, hot-pressed  degradation at  diboride-based composites show a significant loss of strength below 15001C (s1500 o 0.5sRT),115,118,121,123,141 which is well below temperatures expected in proposed applications (420001C). Typically, oxide  form grain boundary  impurities  Fig. 7.  Room temperature flexure  strength as a function of volume  percent SiC and MoSi2 in zirconium diboride -based ceramics produced by hot pressing at 19001C.7,14  phases or become localized at duce strength.121  triple grain junctions, which re In contrast, SPS has been used to produce HfB2-SiC and HfB2-MoSi2 ceramics that maintain their room temperature  \\x0c', '1356  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  appears to produce materials with superior strengths at elevated temperature (s15001C is equivalent to sRT), which merits further investigation. To utilize these materials in the proposed appli cations, the mechanical response must be understood for temperatures from 15001 to 25001C (or higher).  (2)  Elastic Modulus, Hardness, and Toughness  Room temperature elastic modulus (E), Vickers hardness (HV), and fracture toughness values for ZrB2 and HfB2, pure or with SiC or MoSi2 additions, are summarized in Table VII. Elastic modulus for ZrB2 (B489 GPa)7 from recent studies is consistent with values from historic studies using similar sonic resonance techniques on polycrystalline samples.161 Elastic modulus values for HfB2 range from 480 to 510 GPa,8,162 although a value of B445 GPa was reported for porous HfB2.12 Information on elastic constants as a function of temperature has been reported.162 In general, elastic moduli of diboride composites scales with the volume fraction of additive (ESiC \\x18 475 GPa162 and EMoSi2 \\x18 440 GPa163). Hardness shows a similar trend, generally following mixing  rules in relation to the amount and type of phases included in  the particulate composites. Reported hardness values were 21- for polycrystalline ZrB2 and B28 GPa9 23 GPa2,7 for HfB2. Therefore, additions of SiC (HV B28 GPa164) to ZrB2 result in a slight increase in hardness,7 while SiC additions to HfB2 result in little or no change in hardness.115 On the other hand, MoSi2 has a low hardness (B9 GPa165) and its addition decreases the hardness of either ZrB2 113 or HfB2 114-based ceramics. The fracture toughness of ZrB2 and HfB2, with and without is generally in the range of 3.5-4.5 MPa \\x01 m1/2. Fracadditives, ture toughness values (Table VII) have been obtained using a  variety of measurement  techniques,  so direct comparisons are  difficult. In a systematic study of the effect of additive content, Chamberlain et al.7 reported that fracture toughness increased from 3.5 MPa \\x01 m1/2 to 5.3 MPa \\x01 m1/2 for pure ZrB2 for ZrB2 with 30 vol% SiC. Specimens exhibited crack deflection and  crack bridging (Fig. 11). The ZrB2 grains transgranular manner and cracks deflected at or near ZrB2-SiC interfaces, leaving SiC particles in the wake of the advancing cracks.153 This result is consistent with the residual stresses pre typically failed in a  dicted at ZrB2-SiC interfaces due to the mismatch in thermal and mechanical properties between the dispersed SiC particu lates and the ZrB2 matrix (discussed earlier). Further increases in fracture toughness for diboride-based composites will likely  require second phase additions with a higher aspect ratio (e.g., SiC platelets or rods/whiskers)166,167 or the fabrication of laminate-type architectures.168,169  (3)  Thermal Stress Resistance  The use of diboride-based ceramics  in applications  such as  rocket nozzles and leading edges may be  limited by thermal  stress resistance (TSR), which is poor for most ceramics. Ther mal stresses, which develop due to the rapid heating or cooling,  are controlled by the thermal and mechanical properties of the material.170,171 Although five thermal shock parameters are commonly used to describe TSR (R, R0 , R0 0 , R0 0 0 , and R0 0 0 0 ),172 the R parameter (Eq. (13)) is typically used as the worst case  scenario where heating or cooling occurs under conditions of  infinite heat transfer:  R ¼ DTmax ¼ sð1 \\x00 nÞ aE  ðin units of  \\x0eCÞ  (13)  temperature  change  is  the maximum allowable  where DTmax under instantaneous heating or cooling that will not initiate fracture, s is the strength, a is the CTE, n is Poisson’s ratio, and E is Young’s modulus. The R0 and R0 0 parameters are the product of R and the thermal conductivity (R0 in W/m) and thermal diffusivity (R0 0 , in cm2 \\x01 1C/s). The R, R0 , and R0 0 parameters predict resistance to crack initiation and are maximized by increasing strength and decreasing Young’s modulus. The R0 0 0 and  Fig. 9.  Room temperature flexure strength as a function of the average  zirconium diboride (ZrB2) grain and SiC particulate size in ZrB2-30 vol% SiC ceramics (data obtained from Rezaie et al.153 and Zhu et al.156).  The linear curve fit  indicates a strong correlation between strength and  the size of the SiC particulates.  flexure strength up to 15001C.114,115,148 The electrical discharge  used to heat the specimen may promote breakdown of the non conductive oxides on the surface of the particles, perhaps removing them before sintering,160 which could explain the im proved high-temperature mechanical performance reported in  recent SPS studies. The HfB2-based ceramics produced by SPS appear to contain fewer oxide impurities (Si-Hf-C-O) than had  been found in ceramics produced from the same powders using conventional HP.115,148  Figure 10 is a plot of strength as a function of test temperature for HfB2-SiC produced in historic studies,100 and three recent HfB2-based ceramics produced by HP,149 RHP,141 and SPS.115 Strengths were obtained in four-point bending on similar-sized test bars (B4.9-5 mm2 in cross-sectional area). How ever, more strength testing is needed at elevated temperatures.  Reporting strength at two temperatures does not provide com pelling evidence for the effect of temperature (the dashed lines in  Fig. 10 are only included to show general trends). Second, SPS  Fig. 10. Flexure strength as a function of test temperature for hafnium diboride (HfB2) containing 20 vol% SiC produced in the 1960s (100) and three recent HfB2-based ceramics with 22.1-30 vol% SiC additions produced by hot pressing (HP) (Monteverde149), RHP (Monteverde141), and SPS (Bellosi et al.115). The dashed lines for the recent HP, reactive HP ,  and Spark plasma sintering studies are only meant to draw the eye to the  suggested trend in the data.  \\x0c', 'R0 0 0 0 parameters predict crack propagation resistance and can be increased by increasing Young’s modulus and decreasing strength.173 However, resistance to crack initiation is arguably  more important for the applications proposed for diborides be cause of the high heating rates and steady-state heat fluxes. In historic studies, R parameters of 601-1201C were calculated for ZrB2 and HfB2 from room temperature up to B12001C, while the R0 parameters were 2600-4500 W/m.99 The diborides offered a clear advantage over other structural ceramics avail able at the time (BeO, Al2O3, SiC, ZrC, and MgO), which was due, largely, to their strength at elevated temperatures. The  beneficial  role of SiC or SiC1C additions was quantified in improved R0 performance, with ZrB2-based ceramics terms of achieving B9600 W/m under steady-state heat flux.100,174  No subsequent studies have been published that support the  potential advantages of diborides or their composites in terms  of their TSR, although a recent study at UMR has focused on  experimental measurement  of R using water  quench  tests  (see Sidebar 1). The results  show that modern diborides have  R parameters that are three to four times higher than historical values due to increased strength. Even so, the measured DTmax  (B4001C) is well below what may be encountered in proposed  applications.  One method for improving TSR is to tailor structure on mul tiple length scales to produce architectures that are engineered to  enhance TSR while maintaining load-bearing capability such as fibrous monoliths.177 Koh measured an R of 10001C for monolithic Si3N4, and 14001C for a fibrous monolithic ceramic consisting of Si3N4 cells and BN cell boundaries. The improved performance was attributed to increased resistance to crack propagation (increased R0 0 0 0 parameter). Research at UMR on fibrous monoliths with ZrB2-SiC cells and C-ZrB2 cell boundaries showed an improvement of B10001C in the R parameter  compared with conventional materials. However, analysis did  not necessarily support the same mechanism for improved per formance in ZrB2-based materials.  VI.  Oxidation  Many proposed applications for ZrB2 and HfB2 involve oxidizing conditions and elevated temperatures. While other reactive  environments (e.g., molten metals, propellants from rocket mo tors) and erosive conditions are also encountered,  this section  focuses on static oxidation. The role of additives on oxidation  behavior is also discussed. In addition to conventional furnace  oxidation and thermal gravimetry (TG) in flowing air, some re sults from arc heater testing are reviewed.  (1)  Pure Diborides  The diborides undergo stoichiometric oxidation when exposed to air at elevated temperatures178:  ZrB2 ðcÞ þ 5 2  O2 ðgÞ ! ZrO2 ðcÞ þ B2O3 ðl Þ  (14)  HfB2 ðcÞ þ 5 2  O2 ðgÞ ! HfO2 ðcÞ þ B2O3 ðl Þ  (15)  Both reactions are favorable at all temperatures with rxn ¼ \\x001977 þ 0:361T (kJ) for Reaction (14) and \\x002003 1 DGo 0.374 (kJ) for Reaction (15). Analysis by TG shows negligible mass gain below B7001C (Fig. 12). A similar response has been  observed for HfB2, except that its mass gain has been reported to be significantly lower than ZrB2 at all temperatures.179 At below B11001C, ZrO2 (or HfO2)  temperatures  and  B2O3  Table VII.  Room Temperature Elastic Modulus, Hardness, and Fracture Toughness Values for ZrB2and HfB2-Based Ceramics  Composition (vol%)  Modulus (GPa)  Hardness (GPa)  Toughness (MPa-m1/2)  Reference  ZrB2 ZrB2110 SiC ZrB2120 SiC ZrB2130 SiC ZrB2110 MoSi2 ZrB2115 MoSi2 ZrB2120 MoSi2 ZrB2130 MoSi2 ZrB2115 a-SiC14.5 ZrN ZrB2135 HfB2110 a-SiC14.5 ZrN ZrB2137.5 HfB2119.5 a-SiC13HfN HfB2 HfB2120 SiC HfB2130 SiC HfB2115 MoSi2 HfB2120 SiC13HfN HfB2122.1 SiC15.9 HfC HfB2130 SiC12 TaSi2  489-493  21-23  3.5-4.2  2,7,100  450-507  24  4.1-4.8  7,118  466-531  24  4.4  7,100  484  24  5.3  7  516  20.4  4.1  113  479  14.9-16.2  3.3-4.4  115  489-523  16-18.5  2.3-3  113,135  494  17.7  4  113  467  15.6  5  123  494  16.7  4.8  123  497  22  —  121  445  —  —  12  544  17.25  4.15  116  512  26  3.9  115,148  519-530  15.7-20.9  3.77-3.82  114  506  22.3  3.8  115  520  19  —  141  489-506  —  3.6-4.65  149  Fig. 11. A thermally etched cross section of zirconium diboride (ZrB2)- 30 vol% SiC hot pressed at 20501C for 90 min in argon. The image  shows a Vickers  indentation crack path with inset arrows  indicating  predominantly transgranular fracture for the ZrB2 grains and crack deflection near the ZrB2-SiC interfaces.  May 2007  Diborides of Zirconium and Hafnium  1357  \\x0c', '1358  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  Sidebar 1.  Thermal Shock Behavior of ZrB2-Based Ceramics  Table S1. Material Properties and Thermal Shock Parameters for ZrB2 and ZrB2-30 vol% SiC  Material property  Density (g/cm3)  Flexure strength (MPa)  Young’s modulus (GPa)  Assumed Poisson’s ratio Calculated R (1C) Experimental R (1C)  Stress reduction factor  ‘‘Refined’’ R  ZrB2  6.27  568  507  0.16  140 B385 0.364  ZrB2-30 vol% SiC  5.33  823  503  0.16  200 B395 0.364  (calculated)  (assumed)  140  144  A major challenge for ceramics in thermostructural applica tions is their poor resistance to thermal shock. The primary  failure mechanism for pristine monolithic ZrB2 during thermal shock is crack initiation when the stresses imposed by a  thermal  gradient  exceed the  strength of  the material. The  maximum temperature  change  (i.e.,  the R thermal  shock  parameter)  can be  calculated using Eq.  (S1) and measured  properties.  ð  R ¼ s 1 \\x00 n E a  Þ  c  ðS1Þ  where r is the fracture strength, n is Poisson’s ratio, E is the Young modulus, a is the coefficient of thermal expansion, and W is a stress reduction factor. Based on the properties listed in Table S1, the calculated R parameters for ZrB2 and ZrB2- 30% SiC were 1401 and 2001C, respectively. The values can be compared with a calculated R-parameter of B811C for monolithic ZrB2 from historic studies where the fracture strengths were considerably lower (B300 MPa vs. B823 MPa for the current study). The lower strength of the historic material was due to larger grain sizes from higher HP temperatures.175 The experimental thermal shock temperature (R) has been  on  size B  vol% SiC using water  determined for ZrB2 and ZrB2-30 quench thermal shock testing (ASTM C1525) (3 mm \\x02 4 mm \\x02 45 mm) test bars. The thermal shock temperatures for ZrB2 and ZrB2-30% SiC were 3851 and 3951C (Figure S1). A stress reduction factor (W) of 0.36 was required to equate the experimental R for ZrB2 to the calculated value. If the same W were assumed for ZrB2-SiC, the ‘‘refined’’ R for ZrB2-SiC would be 1441C, which is about the same as ZrB2, compared with a calculated value of 1901C. The discrepancy between the calculated and refined R values for ZrB2-SiC may be due to microcracking within the particulate composite, which has been reported for a similar diboride material.176 In addition to conventional materials, R values were meas ured for ZrB2-based fibrous monoliths. As shown in Figure S2, a fibrous monolith was fabricated with cells composed of  ZrB2-30 vol% SiC and cell boundaries composed of graphite- 15 vol% ZrB2. In contrast to the conventional materials, the fibrous monolithic material had an experimental R parameter of B14001C (Figure S1), which is B350% greater than ZrB2 or ZrB2-SiC. The strength of the fibrous monolithic material remained constant for T values up to B14001C, which represents a marked improvement over the conventional ceramics that lost strength for T values of B4001C. Previously, Si3N4/ BN fibrous monoliths have also demonstrated a significant  improvement in thermal shock behavior, although the Si3N4/ BN material had a gradual loss in strength as quench tem perature increased that was not observed in the ZrB2-based fibrous monolith.177 The improved thermal shock resistance of the ZrB2-based fibrous monolith compared with conventional ZrB2 and ZrB2-SiC has been attributed to the fibrous monolithic structure, which limits the initiation and propaga tion of cracks to the cell boundaries, allowing the strong cells  to remain pristine.  form a continuous layer that provides passive oxidation protection.179-184 Analysis has concluded that the rate limiting step for  oxidation is the transport of oxygen through B2O3, which results in parabolic (diffusion-limited growth) kinetics for mass gain  and the oxide  layer  thickness, whether  the  reaction rate was  measured by mass gain, oxygen uptake from the atmosphere, or reaction layer thickness.180,185 Between B11001 and B14001C, the weight change reflects a  combination of mass  loss due to B2O3 evaporation and mass  Fig. S1.  Flexure strength as a function of change in temperature for  water quench thermal testing (ASTM C1525) for ZrB2, ZrB2-30 vol% SiC and ZrB2—30 vol% SiC/graphite—15 vol% ZrB2 fibrous monoliths.  Fig. S2.  Scanning electron microscopy image  showing the macro structure of a ZrB2-SiC/graphite-ZrB2 fibrous monolith. The architecture consists of ZrB2-30 vol% SiC cells with a graphite-15 vol% ZrB2 cell boundary.  gain due to the formation of condensed oxides.182,185 Specimens  continue  to gain mass as  formed is greater  (ZrO2 or HfO2) than the mass of diboride reacted plus the  the mass of oxide  mass of B2O3 lost. As B2O3 evaporates, a porous ZrO2, layer remains, although a small amount of B2O3 may be retained.143 Above B14001C, the oxide layer is not protective and rapid, linear mass gain kinetics have been reported.180 The TG curve  for ZrB2 rapid mass  in Fig. 12 has an increasing slope,  reflecting more  gain at higher  temperatures.  In addition to the  \\x0c', 'May 2007  Diborides of Zirconium and Hafnium  1359  Below B11001C, the addition of SiC does not alter the oxidation behavior of the diborides.182 In this temperature regime, the  oxidation rate of SiC is several orders of magnitude slower than that of the diborides.190 Consequently, the oxide scale on ZrB2- SiC below 11001C consists of ZrO2 and B2O3, as it did for pure ZrB2.189 Above B11001C, two factors affect oxidation. First, the rate of SiC oxidation increases and the SiC particles are  converted to SiO2 plus CO or CO2. Second, the rate of B2O3 evaporation becomes significant. As shown in Fig. 12, ZrB2-SiC shows a mass loss between 12001 and 13001C due to B2O3 evaporation. The silica-rich layer provides protective behavior, which  results in mass gain with parabolic kinetics from room temperature up to at least 16001C.21 Analysis of the outer oxide layer formed at 15001C has detected less than 1 wt% B,  indicating  that nearly all of the B2O3 has evaporated by this temperature.192 The addition of SiC not only extends the temperature  range of protective behavior, but  it also imparts the ability to  rapidly regain protective behavior after  the loss of protection  due to excessive temperature, removal of oxide by shear forces,  or other causes.  Both ZrB2-SiC and HfB2-SiC ceramics exhibit passive oxidation protection with parabolic mass gain kinetics over a wide  range of  temperatures. The oxidation rate is controlled by dif fusion of oxygen through the outer oxide scale. Most authors  also report the formation of a layer containing both ZrO2/HfO2 and SiO2 beneath the outer layer. Often, this region is thin compared with the outer silica-rich scale, but thicker layers have cycling.153 Below the ZrO2-SiO2 (HfO2-SiO2) layer, some authors report the formation of a porous region from which SiC has been depleted.153,190,193-195 This region has been reported to form at temperatures of 15001C or  been reported after  thermal  above and it contains ZrO2 (HfO2), ZrB2 (HfB2), or both. In dry air, thermodynamic modeling suggested that SiC was removed  by active oxidation at the low oxygen partial pressures that are scale.196 Subsequent  to exist under  experimental  thought  the  studies confirmed that a porous ZrO2 layer was formed when ZrB2-SiC was oxidized at 15001C in an oxygen partial pressure of B10\\x0010 Pa. Other compounds such as SiC, SiO2, or B2O3 were not detected in the layer.197 such as MoSi2,113,198 Besides SiC, additives tantalum compounds,119,199 ZrSi2,195 and other diborides21,119,200 improve the  Fig. 12. Weight  change measured  by  thermal  gravimetric  analysis  (TGA) for (a) zirconium diboride (ZrB2) and (b) ZrB2 containing 20 vol% SiC heated at 101C/min to 15001C in air.  experimental studies,  the oxidation behavior of ZrB2 has been described using various thermodynamic models that focus on oxidation in air at 15001C or below.21,186,187  (2)  The Effect of Additives  Additives that alter the composition of the oxide scale improve  the oxidation resistance of ZrB2 and HfB2. By far, common additive is SiC, which reduces the oxidation rate for both ZrB2 and HfB2 by forming a silica-rich scale.177,188-191 Analysis by TG (Fig. 12) shows that ZrB2-SiC had a normalized mass gain of B0.02 mg/mm2 when heated to 15001C in air compared with a mass gain of B0.12 mg/mm2 for ZrB2. A SiC content of 20 vol% has been studied extensively based on his the most  toric studies indicating that this composition had the best combination of oxidation resistance and mechanical behavior.107  Fig. 13.  Cross-sectional scanning electron microscopy image of zirconium diboride (ZrB2)-SiC after arc heater testing at 350 W/cm2 showing (a) the mixed oxide layer, (b) the SiC-depleted layer, and (c) the underlying ZrB2-SiC.  \\x0c', 'oxidation resistance of diborides either alone or in combination  with SiC. In particular, the addition of Ta compounds has been  shown to improve the oxidation resistance of ZrB2- SiC.119,195,201 These additions alter the composition of the outer  glassy layer, which can lead to phase separation of  the liquid/  glass due to the high cation field strength of the transition metals.119,201 In the case of Ta additions, the phase-separated glassy  layers  resulted in reduced oxidation rates  in TG studies with  ZrB2-SiC containing TaB2 having the lowest oxidation rate of any of the materials evaluated.201  (3)  Arc Heater Testing  Although less common than TG or furnace oxidation studies,  arc heater  z  testing is widely recognized as more representative of  the environment that will be encountered in hypersonic flight or  atmospheric re-entry because of its high heat fluxes, low pressure  dissociated atmosphere (i.e., O instead of O2), and high gas velocities.116,196,195,201 Despite the drastic differences in the  experimental conditions, the response of boride-based ceramics  to arc heater testing appears to be similar to conventional studies. After testing at a heat flux of 350 W/cm2 for 10 min (surface modified temperature B18001C), ZrB2-SiC formed a surface layer of ZrO2, a mixed oxide layer, and a SiC-depleted layer (Fig. 13).196 HfB2-SiC showed a similar response under the same test conditions.116 A solid solution modified composition,  (Zr0.9,Cr0.1)B2-SiC, W/cm2 showed a reduced erosion rate  tested by NSWCCD at a heat flux of 326  compared with other  materials.  For  both  ZrB2-SiC and HfB2-SiC, (ZrO2 or HfO2) that formed on contained a small volume fraction of holes  the  dense  oxide  layer  the  surface  thought  to allow  gaseous  reaction products  (i.e., CO, CO2, boron and silicon oxides) to escape without disrupting large sections of the scale.116,196 Based on these arc heater  evaluations, ZrB2-SiC and HfB2-SiC exhibit protective behavior under a wide range of testing conditions.  VII.  Summary  Historic and modern studies of the crystal chemistry, synthesis,  densification, microstructure, mechanical properties, and oxida tion of ZrB2 and HfB2 ceramics were reviewed. Zirconium and hafnium diborides have the AlB2 structure (P6/mmm) that is composed of alternating planes of hexagonally close-packed Zr  atoms and graphite-like B atoms. These compounds exhibit par tial or complete solubility with other transition metal diborides,  which allows properties such as hardness and CTE to be con trolled through the application of crystal chemical principles.  Commercial diboride powders, with particle sizes as fine as 2 mm,  are  commonly  synthesized  by  carbothermal  reduction.  Nano-crystalline powders can be synthesized using solution pre cursors, exothermic reactions, or pre-ceramic polymers. While  HP is  the dominant densification technique,  recent advances  have led to pressureless densification processes  through solid state or liquid-phase sintering. Generally, the strength of the di borides increases as grain size decreases with strengths as high as B500 MPa reported for materials with grain sizes of B5 mm.  The addition of second phases such as SiC can increase strength  to over 1 GPa by further reducing grain size and promoting fa vorable residual thermal stresses. Further, SPS has been used to  produce diboride-based materials that retain their room temleast 15001C. Progress has also been  perature  strength to at  made with regard to understanding the behavior of diboride based ceramics in static oxidation as well as arc heater testing,  which more accurately reproduces  environments  encountered  during hypersonic flight and atmospheric re-entry.  Although significant gains have been made,  several critical  areas have been identified for future research. In particular, re search is needed to understand oxidation mechanisms and fun damental structure-property relations at elevated temperatures,  if these materials are to be used in hypersonic, atmospheric re entry, or rocket propulsion applications.  Acknowledgments  The effort to write this paper was catalyzed by the ‘‘NSF-AFOSR Joint Work shop on Future Ultra High Temperature Materials’’ held at the National Science  Foundation in January 2004. Dr. Lynnette Madsen, Ceramics Program Director  at NSF, and Dr. Joan Fuller, Program Manager for Ceramics and Non-Metallic  Materials at AFOSR, provided the vision for the workshop, which was funded on  NSF grant DMR-0403004. G. E. H. and W. G. F. gratefully acknowledge con tributions of current and former members of the UMR UHTC research group.  References  1A. E. McHale (ed.) Data Collected from Phase Diagrams for Ceramists, Vol. X.  American Ceramic Society, Westerville, OH, 1994. 2R. Telle, L. S. Sigl, and K. 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Dallek,  ‘‘Oxidation of  ZrB2 Ceramics Modified With SiC and Group IV-VI Transition Metal Borides’’; pp. 144-53 in High Temperature Corrosion and Materials Chemistry III, Edited  by M. McNallan, and E. Opila. The Electrochemical Society, Pennington, NJ,  2001. 201J. D. Bull, D. J. Rasky, and J. C. Karika,  ‘‘Stability Characterization of  Diboride Composites Under High Velocity Atmospheric Flight Conditions’’;  pp. T1092-1106 in Advanced Materials: Meeting Economic Challenges, Proceed ings of the 24th International SAMPE Technical Conference, October 20-22, 1992,  Toronto, Canada, Edited by M. Rosenow, T. J. Reinhart, F. H. Froes, and R. A.  Cull. Society for the Advancement of Material and Process Engineering, Covina,  CA.  &  Bill  Fahrenholtz  is  an Associate  Professor  of Ceramic Engineering  at the University of Missouri-Rolla.  He  earned B.S. and M.S. degrees  in  Ceramic  Engineering  at  the  University  of  Illinois  at Urbana Champaign in 1987 and 1989,  re spectively. He completed his Ph.D.  in Chemical Engineering at the Uni versity  of New Mexico  in  1992.  From 1993 to 1999, Bill was a re search assistant professor at UNM.  Since 1999, he has been at UMR where he has received several  campus-wide faculty excellence and teaching awards as well as a  CAREER award from the National Science Foundation. He  teaches  undergraduate  and  graduate  courses  on  thermody namics as well as a required sophomore  level  laboratory on  traditional ceramics. His research focuses on the processing and  characterization of ceramics and ceramic-metal composites. He  has current projects related to ultra-high temperature ceramics  as well as  the use of cerium oxide coatings  for  the corrosion  protection of high-strength aluminum alloys. He has published  over 50 papers on his work.  Greg Hilmas is an Associate Profes sor of Ceramic Engineering at  the  University  of Missouri-Rolla. He  earned B.S.  in Materials  Science  and Engineering at  the University  of Minnesota, his M.S. degree  in  Ceramic Engineering at The Ohio  State University, and his Ph.D.  in  Materials Science and Engineering  from the University of Michigan in  1986, 1988, and 1993,  respectively.  From 1993  to 1994, Greg was  a  post-doctoral  researcher at  the University of Michigan. From  1995 to 1997, Greg was the Composite Materials Manager at  Advanced Ceramics Research in Tucson, AZ. Since 1998, he has  May 2007  Diborides of Zirconium and Hafnium  1363  \\x0c', 'been at UMR where he has  received 10 campus-wide faculty  excellence and teaching awards. In addition, Greg has received  an R&D 100 award for  the development of fibrous monolith  materials  for  rock drill bit  inserts  for  the petroleum drilling  industry. He  teaches undergraduate and graduate  courses on  composite materials and the thermal and mechanical properties  of ceramics. His research focuses on the processing, microstruc ture and properties of ceramics and ceramic composites. He has  current projects related to ultra-high temperature ceramics; high  voltage, high-energy density capacitors; and rapid prototyping  of ceramics. He has published over 50 papers on his work.  Inna Talmy is  the Senior Research  Ceramist and Group Leader of  the  Ceramic  Science  and  Technology  Group at NSWCCD. She joined the  Laboratory  in  1983.  Inna  received  both  her M.S.  (1957)  and  Ph.D.  (1965)  degrees  in Ceramic  Science  and Engineering from the Mendeleev  Institute  of  Chemical  Technology  in Moscow, Russia. Previously,  she  worked at the Institutes of Chemical  Technology in Moscow and Prague,  Czechoslovakia. Her primary  research efforts  are  in dielectric  ceramics,  superconductors,  non-oxide  structural  ceramics,  and  ceramic-matrix  composites.  Inna  directed  the  development  of  celsian and phosphate ceramics as candidates for next generation  tactical missile  radomes. Currently, her  research is  focused on  UHTC materials (such as ZrB2-SiC and cermets) with improved oxidation and thermal stress resistance for hypersonic and strategic  missile applications. Talmy’s work has resulted in over 80 publi cations and 20 patents. Her work has been recognized with several  awards  from the U.S. Navy  including  a Meritorious Civilian  Award, the Adolphus Dahlgren Award, a Science and Technology  Excellence Award, and a NAVSEA Scientist of the Year Award.  James Zaykoski  is a materials scien tist in the Ceramic Science and Tech nology Group at NSWCCD, which  he  joined 1985. He  received a B.S.  (1985) in Materials Engineering from  Wilkes University,  and  an M.S.  (1988) and a Ph.D. (1994) in Materi als Engineering from the University  of  Maryland.  His  research  at  NSWCCD has focused on optimized  processing of carbon-carbon compo sites, toughening of ceramics, model ing of polymorphic phase transformations and gaseous diffusion,  developing oxidation protection coatings and high-temperature  ceramic-based adhesives, and developing non-oxide ceramics for  structural applications  in oxidizing environment. His work has  resulted in 20 publications and nine patents.  Jim received the  NSWC Young Professional of the Year Award in recognition of  his accomplishments.  1364  Journal of the American Ceramic Society—Fahrenholtz et al.  Vol. 90, No. 5  \\x0c']"
},{
  "_id": 235,
  "PDF": "Resistance to thermal shock and to oxidation of metal diborides-SiC ceramics for aerospace application.pdf",
  "Text": "['Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC  Ceramics for Aerospace Application  Fre´ de´ ric Monteverde  w  CNR-ISTEC, Institute of Science and Technology for Ceramics, National Research Council 48018 Faenza, Italy  Luigi Scatteia  CIRA, Italian Aerospace Research Center, 81043 Capua, Italy  Two SiC-containing metal diborides materials, classified in the  ultra-high-temperature  ceramics  (UHTCs) group, were  fabri cated by hot-pressing. SiC, sinterability apart, promoted resist ance  to  oxidation  of  the  diboride matrices. compositions, oxidized in air at 14501C for 1200 min, had mass gains lower than 5 mg/cm2. Slight deviations from parabolic oxidation kinetics were seen. The resistance to thermal  Both  the  shock (TSR) was studied through the method of the retained flexure strength after water quenching (201C of bath temperature). Experimental data showed that the (ZrB21HfB2)-SiC and the ZrB2-SiC materials retained more than 70% of their initial mean flexure strength for thermal quenchs not exceeding 4751 and 3851C, respectively. Certain key TSR properties (i.e., fracture strength and toughness, elastic modulus, and thermal  expansion coefficient) are very similar for the two compositions.  The observed superior critical thermal shock of the (ZrB21HfB2)-SiC composite was explained in terms of more favorable heat transfer parameters conditions that induce less  severe thermal gradients across the specimens of small dimen(i.e., bars 25 mm \\x02 2.5 mm \\x02 2 mm) during the quench down in water. The experimental TSRs are expected to approach the calculated R values (1961 and 2181C for ZrB21HfB2-SiC and ZrB2-SiC, respectively) as the specimen size increases.  sions  I.  Introduction  TRANSITION metal diborides (MB2) such as ZrB2 and HfB2, commonly referred to as ultra high-temperature ceramics (UHTCs), possess melting temperatures greater than 32001C.  This physical characteristic enables interesting perspectives for temperatures above 18001C, a typical  applications at  tempera ture limit of most structural ceramics. Fields within which ZrB2 has already found utilizations include induction-heating parts,  small crucibles for molten metals, dies for wire-drawing, guides  for cold and hot-rolling of alloys and special metal sheets, and  high-temperature electrodes. Aerospace research in the past dec ade has been focused on UHTCs as candidate materials to in crease  the  heat  resistance  of  structural  thermal  protection  systems (TPS) such as leading edges and nose-cones for a new generation of sharp-shaped hypersonic re-entry vehicles.1-3 In deed, UHTCs  could potentially allow space vehicles  to with stand during 18001C,  atmospheric  re-entry  temperatures  in excess of  that  is widely accepted as  the single-use temperature  limit of current hot structures materials such as SiC-coated C-C composites.4 On-ground arc-jet testing has recently shown that  UHTCs can be considered an effective enabling technology for sharp-body space vehicles.2,5,6 Among the family of materials  classified as UHTCs, ZrB2 has a comparatively low density (i.e., 6.09 g/cm3), and has been addressed as one of the most prom ising candidates for this specific aerospace application.  The manufacturing of dense MB2 required the help of applied pressures and prolonged holds in  compacts has  typically  atmosphere-controlled furnaces at sintering temperatures above 20001C because they own strong covalent bondings and low self-diffusion coefficients.7,8  Instead, a number of  studies has  shown that the addition of SiC to MB2 powders has beneficial effects not only on sinterability,9-11 but also on mechanical properties9-12 and resistance to oxidation.1-3,6,13-16 When MB2 is exposed to air at elevated temperatures for instance, MO2 and B2O3 are formed, while the volatilization of B2O3 above 12001C leads to the formation of a MO2 scale.17-20 SiC containing MB2-based composites exhibit improved resistance to oxidation compared with the monolithic diboride  compacts,  thanks to the formation of a silica-based glassy product which  covers the faces exposed to air and provides an effective barrier  to oxygen transport.  As far as the the resistance to thermal shock (TSR)  is con cerned, despite the significance of this physical property for the  UHTCs group, a lack of data characterizes this subject. Dibor ides are  referred to possess good TSR when compared with  other ceramics because of favorable characteristics like the thermal conductivity (i.e., 50-100 W/m \\x01 K)21 and strengths up to 1 GPa.11 Instead, UHTC compositions in the ZrB2-SiC system reported values in the range of 2501-3501C22,23 that signify TSR lower than that of a structural ceramic like Si3N4.24,25 It is widely shared that the thermal shock behavior is influenced by a  number of mechanical and thermo-physical characteristics like  elastic modulus,  strength,  fracture  toughness,  thermal  expan sion, thermal conductivity, and heat transfer rate (i.e. surface conductance).26 The conditions for crack initiation and propagation have been extensively studied by Hasselman,27 while a  series of parameters was defined to relate thermo-physical and  mechanical properties of the materials to their TSR,  i.e the so called thermal stress fracture resistance parameters. The value of  thermal shock at which the scattering of  the retained strength  data significantly spreads out is typically reported as shock DTC. Measurements of elastic moduli at room temperature provide indications over the  critical  thermal  retained strengths and  extent of damage after  thermal  shock.25,28-30  In addition, dif ferent authors argued that size and geometry of  the specimens  have a substantial impact on the extent of damage induced dur ing thermal shock, verifying that it varies significantly from mato material.25,26,30-34 A lack of quantitative  terial  agreement  between (calculated) TSR parameters and experimental DTC was repeatedly pointed out: the importance over such discrepancies of factors like specimen’s size and/or geometry  and temperature dependence of the heat transfer rate coefficient was assessed.25,26,33,34  P. Becher—contributing editor  w  Author to whom correspondence should be addressed. e-mail: fmonte@istec.cnr.it  Manuscript No. 22061. Received July 27, 2006; approved January 9, 2007.  Journal  J. Am. Ceram. Soc., 90 [4] 1130 - 1138 (2007)  DOI: 10.1111/j.1551-2916.2007.01589.x  r 2007 The American Ceramic Society  1130  \\x0c', 'The purpose of  this study was to measure and compare the  TSR of  two SiC-containing MB2-based ceramics, designed as possible base materials for structural  specifically  aerospace  TPSs.  In addition,  in-air oxidation tests were also performed.  Even though similar heat treatments do not fully reproduce the  real operative re-entry conditions, the results provided precious  complementary informations for the understanding of  the oxi dation behavior. The materials herein examined were already characterized in their radiative and surface catalytic behavior35:  thanks to low catalycity and high emissivity, they showed a good  potential for the selected application.  II.  Experimental Procedure  (1) Materials Processing  Two MB2-SiC mixtures of vol%,  commercial powders, amounts  in  ZrB2 þ 15SiC ðcomposition ZSMÞ ZrB2 þ 15SiC þ 10Hf B2 ðcomposition ZHSMÞ  were ball-milled in absolute ethyl alcohol for 24 h, using highpurity zirconia milling media. A quantity o5% of MoSi2 serving as sintering aid was batched into both the compositions.  Table I  shows  some characteristics of  the raw powders used.  After drying in a rotating evaporator, the powder mixtures were sieved through a mesh screen with 250 mm openings.  The powder mixtures were uniaxially hot-pressed in vacuum  using  an inductively heated graphite die,  lined with a BN sprayed graphitized sheet. Peak temperatures/dwell plied pressures were 18201C/15 min/30 MPa for ZSM, and 19001-19401C/45 min/40 MPa for ZHSM, about 201C/min av times/ap erage heating rate. The temperature was measured by means of  an optical pyrometer focused on the graphite die.  (2) Materials Characterization  (A)  Microstructure:  The bulk density (dB) and the theoretical density (dTH) were evaluated using the Archimedes method (water as immersing medium) and the rule-of-mixture,  respectively. The relative density (RD) was calculated dividing  the bulk density by the theoretical density . The phase compos ition was analyzed with an X-ray diffractometer (XRD, Ni-filCuKa  tered  radiation, model D500,  Siemens, Karlsruhe,  Germany) and a scanning electron microscope  (SEM, model  S360, Leica Cambridge, Cambridge, UK) combined with an en ergy-dispersive X-ray microanalyser (EDX, model INCA Ener gy 300, Oxford Instruments, Abington, UK). Polished sections  of the as-sintered materials were prepared with successively finer diamond-based abrasives ranging from 50 to 0.25 mm. A quali tative estimate of the grain size was made off from SEM images  of fracture surfaces.  (B)  Properties  (a)  Mechanical:  The Young’s modulus (E) and Poisson’s ratio (n) were measured on a 28 mm \\x02 8mm \\x02 0.8 mm plate using the resonance frequency method. Micro-hardness (HV1.0)  was evaluated by a Vickers indenter, using a 9.81 N applied load for 15 s. Flexure strength (s) in a 4-pt. configuration was tested  (model Z050, Zwick/Roell—Ulm, Germany) at room temperature (five specimens) on 25 mm \\x02 2.5 mm \\x02 2 mm chamfered bars using 20 mm and 10 mm as outer and inner span, respec tively, and a cross-head speed of 0.5 mm/min. The surface finish  of the bars, measured using a contact stylus type profiler (model  Talysurf Plus, Rank Taylor Hobson, Leicester, U.K.), was Ra 5 0.1470.01 mm, Ra being the arithmetic mean deviation of the assessed profiles. Fracture toughness (KIc) was measured ture (three specimens) using 25 mm \\x02 2 mm \\x02 2.5 mm bars on through the chevron notched beam method at room temperathe same jig used for the flexure strength (cross-head speed 0.02  mm/min). The bars were notched with an 0.08-mm diamond  saw;  the chevron notch tip depth and the average side length  were about 0.12 and 0.8 times the bar thickness, respectively. The ‘‘slice model’’ equation of Munz et al. 36 was used for the  calculation of KIc.  (b)  Thermo-Physical:  Thermal  expansion up to 13001C  was evaluated using a dilatometer (model DIL 402E, Netzsch Gera¨ tebau GmbH, Selb, Germany) in a stream of argon, 51C/ min heating rate. Thermal diffusivity (DTH) up to 12001C was measured through the laser flash method (model LFA-427,  Netzsch Gera¨ tebau GmbH), and following the  standard EN  821-2 as guideline. Heat capacity (CP) was measured using a modulated differential scanning calorimetry (model MDSC, TA  Instruments, New Castle, DE).  The TSR was studied through the method of the retained flexure strength after water-quenching (201C of bath temperature and 501C of thermal quench step). Chamfered bars 25 mm \\x02 2.5 mm \\x02 2 mm were used for testing. The experimental critical thermal shock DTC was determined using the guidelines outlined by the standard prEN820-3. In particular, DTC was identified using a linear interpolation between points that  first reduce the average flexure strength of the quenched bars by  more than 30% of the mean strength of the as-sintered material.  For  each quench temperature, at  least  three  specimens were  tested. In addition, the variation of the elastic modulus (EM) on the thermal shocked specimens was estimated using the load deflection method, and following the standard prEN 843-2 as  guideline. A cross-head displacement (recorded during the load ing cycle) from 50% of the peak load to 90% of the peak load  was selected in order to minimize strong non-linearities.  (c)  Resistance to Oxidation:  The resistance to oxidation was tested in flowing dry air (20 cm3/min) at 14501C for 1200 min, 301C/min of heating rate and free cooling, using a thermo gravimetric analyser (model STA449 Jupiter, Netzsch Gera¨ tebau GmbH, 10\\x003 mg of accuracy) equipped with a vertically heated Al2O3 chamber. Spacers of zirconia were placed between (dimensions 14 mm \\x02 2.5 mm \\x02 2 mm) and the the specimens Al2O3 holder with minimal contact area. The specimen’s mass was measured before and after exposure. XRD and SEM tech niques were used to analyze the microstructures of the oxidized  samples.  III.  Results  (1)  Microstructure Development  Table II shows processing parameters and some microstructure  features of  the hot-pressed materials. The bulk densities  dB of the as-sintered ZSM and ZHSM compacts were 5.61 and 6.06 g/cm\\x003, respectively. These values of density correspond to relative densities of 99% ZSM and 98.2% ZHSM. The fracture  surfaces observed by SEM displayed the  grain structure of  both the composites, with regularly faceted diboride (maximum size 5 mm). Representative regions of  grains  fracture sur Table I.  Characteristics of the Commercial Powders Used (Source: Companies Datasheet)  Formula  Company  Type  \\x003) Density (g cm  Particle size  Oxygen (wt%)  ZrB2 HfB2 SiC  H.C. Starck (Goslar, Germany)  B  6.09  o2 mm  2  Cerac Inc. (Bad Soden-Salmu¨ nster, Germany)  325 mesh  11.18  1.7 mm 11.6 m2/g 2.8 mm  w  —  H.C. Starck  BF 12  3.19  z  0.9  MoSi2  Aldrich (Milwaukee, WI)  6.26  y  1  w  Fisher size (APS).  z  Specific surface area.  y  Mean value.  April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1131  \\x0c', 'faces  revealed also a high percentage of  transgranular versus  intergranular fracture between the diboride grains in both the  compositions. The XRD analysis of the ZSM compact identified  the main phases ZrB2 and SiC, and an amount of monoclinic zirconia o1 vol%. Instead, in the ZHSM compact, obvious changes in the starting composition took place during sintering.  XRD analyses showed the formation of ZrB2/HfB2 solid solutions, hereafter indicated as (Zr,Hf)B2. Such solid solutions uphold the hexagonal structure of the initial ZrB2, but host at the same time Hf atoms. The presence of (Zr,Hf)B2 solid solutions, which very often appear as shells around ZrB2 cores, is exactly confirmed by the typical configuration of the diboride grains  structured  in  core-shell  (Fig.  1).  The  formation  of  such  (Zr,Hf)B2 s.s. was further assessed through localized EDX analyses on polished surfaces. Furthermore, the polished section  examined by SEM highlighted the SiC particulates distributed  intergranularly within the diboride matrices, in ZSM sometimes in agglomerates (maximum size 5 mm). At the same time, SEM  observations do not provide evidence of residual porosity.  Compared with the ZSM composition, the addition of HfB2 into the ZHSM composition enhanced refractoriness of the in itial powder mixture. In fact, higher peak temperature and long er hold were necessarily applied so far to complete densification: during the hold of 45 min, temperature drifted from 19001 to 19401C. Such conditions led the grain structure of  the ZHSM  composition to increase in the average size more than that of the  ZSM composition (Table  II). Moreover,  in ZHSM a partial  sintering  of  the  intergranular  SiC agglomerates  took  place.  Processing,  densification  behavior,  and microstructure  deve lopment  of  similar  systems  are  described  in  several  earlier  papers.9-11,21,24,37  (2)  Mechanical and Thermo-Physical Properties  Tables III and IV and Figs. 2-4 show the experimental data for  selected properties. Average values of flexure strength at room temperature (sR0) of 795 and 765 MPa for the as-sintered ZSM and ZHSM materials, respectively, result better than for a sim ilar system composed of ZrB2 mixed with 15 vol% SiC and then hot-pressed.37 The co-existence of very hard components like  MB2 and SiC, along with the fine and dense structure, favored high levels of micro-hardness. However, 4.1 MPa \\x01 m1/2 as mean value of fracture toughness in both the materials is typical for this family of (brittle) ceramics.2,9-12,24,37 The behavior of the retained strength (sR) after water-quench is shown in Fig. 2. The sR data of the thermal quenched bars in Fig. 2 indicate that a thermal shock severity DTC as high as 3851 and 4751C were sustained by ZSM and ZHSM, respectively,  before losing more than 30% of their initial mean flexure strength sR0 (Table III). Just for comparison, a ZrB2 material containing 15 vol% SiC (95% relative density and 235 MPa  mean flexure strength), tested using the water-quenching method, reported a TSR of 3501C.23 It should also be noted that the ZSM specimens quenched above DTC have an evident scatter of strength values above and below the reference threshold (i.e.,  70% of the initial mean strength). On the contrary, the ZHSM material seems to follow Hasselman’s theory,27 which predicts a  sharp drop in strength at a critical  thermal shock temperature  difference. The plots in Fig. 3 show the elastic modulus (EM) of a thermal shocked specimen against its sR. A substantial decrease in strength was very often accompanied with a significant re duction in elastic modulus.  In order to confirm that the decrease in flexure strength after  quenching was actually an effect of the thermal shock, four specimens of ZSM were heated up to 5001C (51C/min heating  rate), held at  that  temperature for 15 min to establish thermal  equilibrium, and then left cooling freely inside the furnace instead of quenching them into the water bath at 201C. The re850715 MPa mean71  tained  flexure  strength,  standard  deviation, measured at  room temperature on the heated (and  not-quenched) ZSM specimens, did not show significative difif compared with sR0 of III). This suggests that the decay in flexural strength measured  ferences,  the pristine material  (Table  after quenching is primarily caused by thermal  shock-induced  effects.  The conditions  for  failure in the pristine material during a  thermal quench is usually predicted by setting the maximum  permitted thermally-induced stress equal  to its tensile or bend  strength. Table IV lists  two TSR parameters calculated from  Eqs. (1) and (2)  R ¼ s ð1 \\x00 nÞ=ðaE Þ  (1)  R0 ¼ RKTH  (2)  where s, n, a, E, and KTH represent modulus of rupture, Poisson’s ratio, coefficient of thermal expansion, elastic modulus,  and thermal conductivity, respectively. The comparison of  the  TSR values in Table IV shows that the calculated R values are much lower than the experimental DTC.  Table II.  Processing Parameters and Microstructure Features  of the Hot-Pressed Materials: Peak Temperature (T ), Dwell Time (t), Maximum Applied Pressure (P), Theoretical (dTH) and Bulk Density (dB), Relative Density (RD), and Grain Size (GS)  Sample  Processing parameters  Microstructure features  T (1C)  t (min) P (MPa)  Density  GS (mm)  dTH (g/cm3)  dB (g/cm3) RD (%)  ZSM  1820  15  30  5.67  5.61  99.0  2  ZHSM 1940  45  40  6.17  6.06  98.2  3  Fig. 1.  Polished cross-sections of ZSM and ZHSM (SEM micrographs, secondary electrons); dark features are SiC particles. The core (C)- rim (R)  structure in ZHSM is indicated.  1132  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  \\x0c', 'Discrepancies between calculated R values and observed DTC have been reported previously by many investigators,25,26,33,34  and may indicate that cracks  initiate under  the so-called soft  thermal shock. These conditions imply that  the interior of  the  specimens cools before the surface, resulting in the need for a  stress reduction factor determined using the so-called Biot modulus b ¼ Lh=KTH (L characteristic specimen’s length, KTH thermal conductivity, and h surface heat transfer coefficient be tween specimen and cooling medium). The R parameter cap tures the initiation of thermal shock cracking under hard thermal shock conditions (i.e., values of b>2026): higher values  of R are in favor of greater resistance to fracture initiation during quenching. Differences between R and DTC are particularly evident for water-quenching thermal shock tests, primarily be cause of the dramatic reduction in heat transfer coefficient after  the formation of protective steam bubbles at the water/specimen  interface:  such a phenomenon would mitigate the severity of thermal shock, and result in overestimated DTC values, compared to the expected R parameter.33  The thermal conductivity (KTH) was calculated using the expression KTH ¼ DTHCP r, where DTH, CP, and r represent thermal diffusivity (Fig. 4), heat capacity (Fig. 4), and density, 301-12001C,  respectively. Within  the  interval  the  calculated  data indicate  that ZHSM tends  to dissipate heat  faster  than  ZSM (Table IV). Just for comparison, such values of KTH result lower than that reported for a ZrB2120%SiC, i.e., 103.8 W/ m1C at RT.38 In the present case of small sized specimens and significant KTH (i.e., low values of b), the TSR ranking is more adequately determined by the R0 parameter however, expected that in increasing L (i.e., the minimum heat  (Table IV).  It  is,  transfer  (specimen) dimension or its thickness in the case of the observed DTC values begin to approach the correspondent R parameters (see Table IV), mini standard flexure bar)  mizing the differences between the two tested compositions.  (3)  Resistance to Oxidation  The graphical trends of the specific mass change (w) versus time (t) during the 1200 min exposure at 14501C is plotted in Fig. 5, 3.4370.02 and 4.8870.02 mg/cm2 being the final w values for ZSM and ZHSM, respectively. Offsets equal to 0.4770.02 and 0.470.02 mg/cm2, which account for some oxidation prior the  1200 min hold, were respectively subtracted from raw data of  ZSM and ZHSM. Some data processing was done by calculating the dummy constant KD 5 w2/t points out that during the isothermal exposure obvious para (Fig. 6). Such an exercise  bolic kinetics (i.e., slope of the log KD equal to zero) were never established. In addition, the analysis of the very early stages of  the isothermal hold (see inset in Fig. 5) shows however prevail ing mass loss mechanisms over mass gain mechanisms, which in  turn become typical for the remaining exposure.  The XRD analyses on the surfaces exposed to air  revealed  monoclinic MO2, M 5 Zr for ZSM, and Zr/Hf for ZHSM, while the cross-sections emphasized a structure that consists of an  outermost glassy layer on top of a sub-scale composed of MO2 crystals (partially embedded within the same glass), which ex tends up to the unoxidized bulk (Fig. 7). The glassy layer, whose  composition analysed by SEM-EDX falls in the Si-O system,  adheres to the sub-scale, even though its thickness varies from  few to some tenths of microns. The values of about 50 and 100 mm are grossly representative of the oxide sub-scale thickness for  ZSM and ZHSM, respectively. In addition, EDX analyses as sessed that, in proximity of the interface between oxide sub-scale  and unoxidized bulk (Fig. 8),  sites  formerly occupied by SiC  particulates now appear occupied by a carbon-based solid com pound. Such an oxidation by-product, whose occurrence was already reported,10 was connected to the active oxidation of SiC,  which induces a partial depletion of SiC close to the inner ox idation front facing the virgin bulk. Such an oxidation mecha nism creates  so  far  porosity within  the  sub-scale  formerly  mentioned.  IV.  Discussion  (1)  Microstucture and TSR  Contrary to the well-known limitations connected to the densification of MB2-based materials,5,8-11,39 the obtainment of near full dense compacts was allowed by SiC which substantially en hanced densification of MB2 during hot-pressing. Similarly to powders of TiB2,40,41 MO2 and B2O3 were assumed as the main oxygen carriers upon the surfaces of MB2, M 5 Zr and Hf. Such a contamination by oxygen promotes vapor phase transport  (and thus coarsening) at temperatures below which mass trans fer mechanisms  like  boundary/volume  diffusion, which  are  much more effective than vapor phase for densification,  start  taking place:  the anticipated coarsening decreases  the driving  force for densification at higher temperatures. Densification of  SiC-containing MB2 powder mixtures peratures compared with pure MB2 are deemed to remove the oxide coatings separating MB2 particles from mutual contact.  initiates at  lower  tem 10,42 as  reactions with SiC  As  far as  the TSR evaluated through the retained strength superior DTC in the ZHSM material, compared with that of ZSM material, was at after water-quenching is  concerned,  the  tributed to more favorable heat transfer parameters conditions  Table III. Mechanical Properties of the Hot-Pressed ZSM and ZHSM Materials: Elastic Moduli E (Resonance Frequency Method) and EM0 (Load-Deflection Method), DE 5 (E-EM0)/E, Poisson’s Ratio n, Micro-Hardness HV1.0, Fracture Toughness KIc, and Flexure Strength at Room Temperature rR0  Sample  E  w  (GPa)  EM0  z (GPa)  DE (%)  n  zHV1.0 (GPa)  zKIc (MPaOm)  zsR0 (MPa)  ZSM  48074 50874  44475 47379  7.5  0.12  17.770.4 18.270.5  4.170.05 4.170.75  7957105 765775  ZHSM  6.9  0.128  w  Uncertainity.  z  Mean71 standard deviation.  Thermo-Physical Properties and TSR Parameters (R and R0 ) of the Hot-Pressed ZSM and ZHSM Materials: Linear Table IV. Coefficient of Thermal Expansion a (251-13001C), Thermal Conductivity KTH, and Critical Thermal Shock DTC  Sample  \\x006/1C) a (10  KTH(W/m \\x01 1C)  TSR parameters  301C  5001C  10001C  12001C  w  R (1C)  w  R0  (kW/m)  DTC (1C)  301C  5001C  ZSM  6.68  62.5  64.5  65.1  65.2  218  13.6  14.0  385  ZHSM  6.74  79.9  83.5  84.2  85.0  196  15.7  16.4  475  w  Calculated using sR0 in Eq. (1). TSR, resistance to thermal shock.  April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1133  \\x0c', '1134  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  Fig. 2. Retained flexure strength (sR) vs. thermal shock (T) of ZSM and ZHSM in water bath at 201C; full tested. Horizontal dashed segments: 70% sR0 of ZSM and ZSM and ZHSM (see Table III).  lines connect mean sR values at each T  that induce less severe thermal gradients across the specimen during the quench down test in a 201C water-bath. Being all the  tions, and are not solely dependent on materials characteristics.  Furthermore,  the determination of adequate  strength values,  specimens  tested of  the same (small) dimensions,  the different  such as the Weibull lower limiting strength,  is known to require  KTH (Table IV) was deemed to modify the actual Biot moduli of the two composites during quenching. It can at first be noticed to strength (s) ratios repre large fracture toughness (KIc) sent a favorable material’s characteristic for minimizing the ex that  tent of the crack propagation, but stands in direct contrast with  the requirement of high strain-to-failure necessary for preventing crack initiation.26,27,43 Hence, the combination of large E and s values which pertains the present materials may induce  deleterious damage resulting from fracture if it does occur upon severe thermal shocking, i.e., above DTC. It was in fact observed that, for thermal quench above DTC, a ZSM specimen spalled into a number of pieces, i.e., it had no measurable retained  strength. This situation very likely arose because of the sudden  release of  the large thermal strain-energy required to initiate a  crack from very small flaws in the material. The observed DTC for ZSM and ZHSM, respectively, seem to contradict the expected ranking based on the calculated R pa rameters. Lewis has  stressed the  importance of  selecting the  most appropriate value for strength so that meaningful rankings can be settled.34 In addition, Becher et al. argued that, as does  extensive testing. On the other hand, when extensive strength  data from appropriately sized specimens are unavoidable (this is  the present case), reasonable success can be obtained by taking a  characteristic strength as the mean value less two or three standard deviations.45 Calculated R values and experimental DTC under regimes of soft thermal quenchs have been repeatedly correlated through empirical expressions DTC 5 f(b) \\x01 R which try to approximate the thermal stress solutions for specific sample geometries of rod, plate, or bar.25,26,30,33,34 Also for the pres ent study, an empirical approach was adopted (damping factor f(b) 5 11B/b, 0o1/f(b)o1, b 5 Lh/KTH, h surface heat transfer coefficient, L minimum specimen dimension, KTH thermal conductivity, and B shape factor). According to the Lewis criterion,45 the mean strengths sR0 less three standard deviations (named sZSM and sZHSM in Fig. 9) were used for calculating corrected R parameters for ZSM (RZSM) and ZHSM (RZHSM), respectively. By setting L 5 2 mm, B 5 4, and KTH 5 83.4 W/ m \\x01 K at 4751C for ZHSM, an effective h of 67 kW/m2 \\x01 1C was inferred, being bD1.6 in correspondence of DTC equal to 4751C. This guess for the h parameter, though merely practical for such  the size dependence of the strength of ceramics, the temperature  a comparison,  falls anyhow very close to the typical  range of  dependence of thermal conductivity and surface heat transfer coefficient varies with temperature for different materials.33 Un fortunately,  surface heat  transfer  coefficients usually are not  known precisely and have been shown, moreover, to vary over  three orders of magnitude within the typical temperature range of water-quench tests. Therefore the actual DTC values for essentially all materials are function of the heat-transfer condi values for ceramics during water-quenching thermal shock tests,34,44 in particular 50 kW/m2 \\x01 1C for TiB2 at 5001C of thermal water-quench.44 Keeping the same values of h and L for the reduced KTH value of 64.2 W/m \\x01 1C at 3851C (that ZSM, means b approaching 2.1) induces more severe thermal gradients  in the quenched specimens, leading to a less effective capacity to withstand thermal stresses (i.e., DTC approaching 3851C). Being  Fig. 3. Retained flexure strength (sR) vs. residual elastic modulus (EM) of bars thermal shocked in water bath at 201C. Horizontal dashed segments: 70% sR0 of ZSM and ZHSM (see Table III).  \\x0c', 'April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1135  Fig. 4.  Heat capacity (CP) and thermal diffusivity (DTH) vs. temperature (T) of ZSM and ZHSM up to 12001C.  the dimensions of the specimens nominally identical, such a difference in DTC can thus be connected to the superior thermal conductivity that in the ZHSM material modulates a more favorable damping factor f(b) through the Biot number. It is also  scatter is associated to concurrent large scatter in retained elastic  moduli: there is no thorough explanation for this. On one hand,  to the  authors’ knowledge,  a  similar behavior  in analogous  systems was not verified. In ZSM material, the concomitant oc expected that, having both the composites levelled characteris currence of large scatter in retained values of strength and elastic  tics  like strength, Poisson’s  ratio, fracture toughness, thermal expansion and elastic modulus, the observed DTC parameters approach the calculated R values of about 2001-2201C in com ponents of larger dimensions.  In Fig. 3, the sharp reduction in elastic modulus EM was very often accompanied with a significant decrease in the retained strength sR. The general agreement between the reduction of flexure strength to coincide with a loss of elastic modulus makes  modulus  (Figs. 2 and 3) may be indicative of a non-uniform  growth in cracks density and/or  size due to sub-critical crack  propagation in the water quench bath. On the other hand, the  quench experiments have some well-recognized limitations. The  TSR tests herein adopted require for instance sample geometries  that are amenable to subsequent mechanical testing, and many  specimens should be tested to assure the significance of the re sults.  In addition,  the  specimens are tested destructively:  this  such combined measurements a useful  tool  to correlate the se implies that initiation sites and sub-critical thermal damage are  verity of the damage introduced during a thermal quench with  difficult  to identify. Thermal  stress gradients across  the small  the capacity to withstand further external loads. Experimental data in Fig. 3 clearly show that sR results more susceptible than EM to the influence of newly formed flaws during thermal shock. According to the Griffith relation among strength s, fracture toughness KIc and flaw size c, sEKIc = growth in size of just a single flaw after thermal shock have more obvious effects on strength than on elastic modulus.46 Evidence  the initiation and  p  ffiffiffi  c ,  of it is in both the compositions, when the elastic modulus declines less or not at all, compared with the corresponding sR values which exceeded an absolute loss of more than 30% sR0 (Table III). Therefore, for apparent drops of the elastic modu sized specimens may therefore contribute to the variability of the retained strength after quenching above DTC due to the probability of a initiation site being positioned in the high stress lo cations.  It  is known in fact  that  the flexure  strength testing  reaches the maximum stress only close to the specimen’s surface.  The combination of the thermo-mechanical characteristics in the  ZHSM material is somewhat more beneficial for a superior TSR  as compared with that of  the ZSM material, especially for the  small specimens herein tested. However, the absence of a large  scatter in retained strength and elastic modulus with increased  thermal shock seen for the ZSM composite offers a real advan lus, the corresponding strength decay is very likely connected to  tage for the ZHSM material. More extensive testing and char a substantial damage initiated and grown during thermal shock  acterizations would assist in identifying the difference in damage  in form of random network of cracks.  mechanisms and the benefits of such behavior.  A significantly larger scatter in retained strength differentiates  ZSM material from ZHSM material (Fig. 3). Also, the strength  (2)  The Resistance to Oxidation  The negative oxidation rate along the very early stages of exposure at 14501C (see inset in Fig. 5) discloses that the oxidation  of MB2  MB2 þ 5=2 O2 ðgÞ ¼ MO2 þ B2O3 ðl Þ  is however dominated by the volatilization of boron oxide  B2O3 ðl Þ ¼ B2O3 ðgÞ  (3)  (4)  from the  specimen’s  surfaces  exposed to the air-atmosphere.  Reactions 3 and 4 describe mechanisms of weight gain and  weight loss, respectively.  With increasing exposure time, the resistances to oxidation of  both the composites, due to SiC which is known to start oxi dizing more slowly than MB2, begin taking advantage from the ongoing formation of a glass coating on the faces exposed to air  (Fig. 7). Generally SiC, according to the reaction  Fig. 5. Specific mass change (w) vs. exposure time (t) during oxidation at 14501C. The inset expands the very early stages of the isothermal hold.  SiC þ 3=2 O2 ðgÞ ¼ SiO2 þ COðgÞ  (5)  \\x0c', '1136  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  Fig. 6.  Dummy parameter KD vs. exposure time (t), KD 5 w2/t.  reacts with oxygen and provides silica. Similarly to other ZrB2- SiC systems,1,3,6,10,13,15,16,24,47 significant amounts of silica com bine with the available B2O3 and yield a borosilicate glass. Such a glassy oxidation product is deemed to lose progressively the  B2O3 component, but is capable of providing better protection against oxidation than an MO2-based scale covered with only boron oxide.6 Although the borosilicate glass behaves as excellent oxidation barrier below 16001C,  it softens, due to the di minished viscosity. The undulating pattern of the external glass  is ascribed to the reduced viscosity.  The marked difference of  the thermo-gravimetric curves  in  Fig. 5 is still a matter of reasoning. The evaluation of the specific  weight changes could have been underestimated because of phe nomena connected to simultaneous gains and losses of mass,  which a previous work on the ZrB2-SiC system firmly asserted.15 The initial faster mass gain rate in ZHSM was attributed to  an inadequate provision of protective glass, compared to ZSM,  coupled with an easier diffusion of oxygen along short-circuited  paths through the forming (Zr,Hf)O2 sub-scale. Next, the reason why, compared to ZHSM, ZSM better resisted against oxida tion remains unclear. Having both the compositions the same  nominal  content of SiC and negligible levels of porosity,  the  core-shell structure of (Zr,Hf)B2 solid solutions in ZHSM supposedly was thought to offer less resistance to the inward dif fusion of oxygen, compared to ZrB2. Even though an oxidation process governed by the oxygen  diffusion through a glass that fits parabolic kinetics  seemed a  logic  expectation, departures  from an ideal parabolic pattern  were instead of not marginal significance (Fig. 6). A number of  factors is indicated responsible of such deviations: the evolution  of volatile products (B2O3, CO), non-steady state behavior due to faster oxidation of MB2 compared with SiC, and the active  Fig. 7. Fracture cross-section of the ZHSM material after 1200 min at 14501C in air  secondary electrons). An oxide sub (SEM micrograph,  scale (100 m thick) underlying the external glass (25 m thick) is indicated.  Fig. 8.  ZHSM material after 1200 min at 14501C in air; SEM micro graph, secondary electrons. SiC particles and carbon-based by products  (arrows) near the oxide sub-scale/unoxidized bulk interface are shown.  oxidation of SiC. The last factor remains being explained thor oughly. The examination through the SEM-EDX technique of  the  sub-oxide  scale/diboride  interface  (i.e  across  the  inner  oxidation front) revealed inclusions containing carbon (Fig. 8).  Being the size, shape, and distribution referable to the SiC par ticulates  in the unoxidized material,  these carbon-based struc tures were associated with an active oxidation mechanism of  SiC. Unlike some authors who claimed the complete volatiliza tion of SiC through active oxidation in SiC-containing MB2 matrices,6,15,24 the detection of C in such solid inclusions led to  consider the following thermodynamic equilibrium  SiC þ SiO2 , 2SiOðgÞ þ C  (6)  as the transition mechanism by which SiC starts experiencing the decomposition.48 The  SiC transforms  active  reason why  in  agreement with equilibrium 6 is not fully understood. Most re cently, Fahrenholtz developed a thermodynamic analysis of the ZrB2-SiC oxidation49: once SiC starts oxidizing according to reaction (7)  SiC þ O2 , SiOðgÞ þ COðgÞ  (7)  CO may further reduce to solid C at the so called sooting limit  (see reaction 8)  2COðgÞ , C þ O2 ðgÞ  (8)  Fig. 9. Biot number b vs. R parameter corrected via damping factor f(b) 5 114/b. According to Lewis criterion,45 sZSM and sZHSM, are sR0 of ZSM and ZHSM, respectively, less 3 standard deviations (see Table III). RZSM and RZHSM are obtained by using sZSM and sZHSM in Eq. (1).  \\x0c', 'In both cases,  it  is  as much as plausible  that  conditions  of high temperature and low oxygen partial pressure inside the the active oxidation of SiC.6,49,50  oxidized scale were met  for  Furthermore, once SiO(g) diffuses outward and encounters a  higher oxygen partial pressure, it would further convert into the  condensed silica phase. Glass pockets inside the oxide sub-scale  but  in the  vicinity of  the  external  glass may be  convincing  evidence of  the phenomenon just described.  It  should also be  considered that  the lack not only of open porosity but also of  easily-oxidizable secondary phases is a favorable condition that  does not offer preferential paths like pores or grain-boundary  channels to the inner transport of oxygen. As a consequence, the  oxygen partial pressure PO2 sufficiently low to induce the active oxidation of SiC. The PO2 parameter in fact is widely recognized as one of the primary  inside the bulk remains at a level  factors controlling the active-to-passive oxidation transition of SiC.48-51 Such an instability of SiC could have not negligible  influence on MB2-SiC systems in oxidizing conditions, dictating limited durability for prolonged services in thermally harsh ap plications.  V.  Summary  The present work explored the resistance to thermal shock and  to oxidation of two ultra-high-temperature MB2115 vol% SiC composites, M 5 Zr and Zr1Hf, brought to full density via hotrather uniform (2-3 mm  pressing. The microstructures were  grain size), with SiC incorporated intergranularly within the di boride matrices. The presence of SiC promoted the resistance to  oxidation of  the diboride matrices  through the coverage of a  silica-based glass which behaved as an effective barrier against  oxidation at high temperature. Both the compositions, oxidized at 14501C for 1200 min, had specific mass gains lower than 5 mg/cm2. Deviations  from parabolic  oxidation  kinetics were  seen, and attributed to evolution and release of volatile prod ucts, to non-steady state behavior based on the faster oxidation  of MB2 compared with SiC, and to active oxidation of SiC. The passive-to-active transition in the oxidation behavior of SiC  makes  the durability of  these SiC-containing diborides-based  composites in long-term thermally severe applications an issue.  The TSR,  tested through the  retained strength after water quench, showed that the ZrB2-SiC and (ZrB21HfB2)-SiC materials retained more than 70% of their initial mean strength (i.e., DTC) once the thermal shock has not exceeded 3851 and 4751C, respectively. On the contrary, the calculated R parameter the ZrB2-SiC and (ZrB21HfB2)-SiC materials were 2181 1961C, respectively. The improved DTC in the explained in terms of more  for  and  (ZrB21HfB2)-SiC was heat transfer parameters conditions  favorable  that establish less critical  thermal gradients across the specimens of reduced dimensions (i.e., bars 25 mm \\x02 2.5 mm \\x02 2 mm) during a quench down test. 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  "PDF": "Resistance to Thermal Shock and to Oxidation of Metal Diborides–SiC Ceramics for Aerospace Application.pdf",
  "Text": "['Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC  Ceramics for Aerospace Application  Fre´ de´ ric Monteverde  w  CNR-ISTEC, Institute of Science and Technology for Ceramics, National Research Council 48018 Faenza, Italy  Luigi Scatteia  CIRA, Italian Aerospace Research Center, 81043 Capua, Italy  Two SiC-containing metal diborides materials, classified in the  ultra-high-temperature  ceramics  (UHTCs) group, were  fabri cated by hot-pressing. SiC, sinterability apart, promoted resist ance  to  oxidation  of  the  diboride matrices. compositions, oxidized in air at 14501C for 1200 min, had mass gains lower than 5 mg/cm2. Slight deviations from parabolic oxidation kinetics were seen. The resistance to thermal  Both  the  shock (TSR) was studied through the method of the retained flexure strength after water quenching (201C of bath temperature). Experimental data showed that the (ZrB21HfB2)-SiC and the ZrB2-SiC materials retained more than 70% of their initial mean flexure strength for thermal quenchs not exceeding 4751 and 3851C, respectively. Certain key TSR properties (i.e., fracture strength and toughness, elastic modulus, and thermal  expansion coefficient) are very similar for the two compositions.  The observed superior critical thermal shock of the (ZrB21HfB2)-SiC composite was explained in terms of more favorable heat transfer parameters conditions that induce less  severe thermal gradients across the specimens of small dimen(i.e., bars 25 mm \\x02 2.5 mm \\x02 2 mm) during the quench down in water. The experimental TSRs are expected to approach the calculated R values (1961 and 2181C for ZrB21HfB2-SiC and ZrB2-SiC, respectively) as the specimen size increases.  sions  I.  Introduction  TRANSITION metal diborides (MB2) such as ZrB2 and HfB2, commonly referred to as ultra high-temperature ceramics (UHTCs), possess melting temperatures greater than 32001C.  This physical characteristic enables interesting perspectives for temperatures above 18001C, a typical  applications at  tempera ture limit of most structural ceramics. Fields within which ZrB2 has already found utilizations include induction-heating parts,  small crucibles for molten metals, dies for wire-drawing, guides  for cold and hot-rolling of alloys and special metal sheets, and  high-temperature electrodes. Aerospace research in the past dec ade has been focused on UHTCs as candidate materials to in crease  the  heat  resistance  of  structural  thermal  protection  systems (TPS) such as leading edges and nose-cones for a new generation of sharp-shaped hypersonic re-entry vehicles.1-3 In deed, UHTCs  could potentially allow space vehicles  to with stand during 18001C,  atmospheric  re-entry  temperatures  in excess of  that  is widely accepted as  the single-use temperature  limit of current hot structures materials such as SiC-coated C-C composites.4 On-ground arc-jet testing has recently shown that  UHTCs can be considered an effective enabling technology for sharp-body space vehicles.2,5,6 Among the family of materials  classified as UHTCs, ZrB2 has a comparatively low density (i.e., 6.09 g/cm3), and has been addressed as one of the most prom ising candidates for this specific aerospace application.  The manufacturing of dense MB2 required the help of applied pressures and prolonged holds in  compacts has  typically  atmosphere-controlled furnaces at sintering temperatures above 20001C because they own strong covalent bondings and low self-diffusion coefficients.7,8  Instead, a number of  studies has  shown that the addition of SiC to MB2 powders has beneficial effects not only on sinterability,9-11 but also on mechanical properties9-12 and resistance to oxidation.1-3,6,13-16 When MB2 is exposed to air at elevated temperatures for instance, MO2 and B2O3 are formed, while the volatilization of B2O3 above 12001C leads to the formation of a MO2 scale.17-20 SiC containing MB2-based composites exhibit improved resistance to oxidation compared with the monolithic diboride  compacts,  thanks to the formation of a silica-based glassy product which  covers the faces exposed to air and provides an effective barrier  to oxygen transport.  As far as the the resistance to thermal shock (TSR)  is con cerned, despite the significance of this physical property for the  UHTCs group, a lack of data characterizes this subject. Dibor ides are  referred to possess good TSR when compared with  other ceramics because of favorable characteristics like the thermal conductivity (i.e., 50-100 W/m \\x01 K)21 and strengths up to 1 GPa.11 Instead, UHTC compositions in the ZrB2-SiC system reported values in the range of 2501-3501C22,23 that signify TSR lower than that of a structural ceramic like Si3N4.24,25 It is widely shared that the thermal shock behavior is influenced by a  number of mechanical and thermo-physical characteristics like  elastic modulus,  strength,  fracture  toughness,  thermal  expan sion, thermal conductivity, and heat transfer rate (i.e. surface conductance).26 The conditions for crack initiation and propagation have been extensively studied by Hasselman,27 while a  series of parameters was defined to relate thermo-physical and  mechanical properties of the materials to their TSR,  i.e the so called thermal stress fracture resistance parameters. The value of  thermal shock at which the scattering of  the retained strength  data significantly spreads out is typically reported as shock DTC. Measurements of elastic moduli at room temperature provide indications over the  critical  thermal  retained strengths and  extent of damage after  thermal  shock.25,28-30  In addition, dif ferent authors argued that size and geometry of  the specimens  have a substantial impact on the extent of damage induced dur ing thermal shock, verifying that it varies significantly from mato material.25,26,30-34 A lack of quantitative  terial  agreement  between (calculated) TSR parameters and experimental DTC was repeatedly pointed out: the importance over such discrepancies of factors like specimen’s size and/or geometry  and temperature dependence of the heat transfer rate coefficient was assessed.25,26,33,34  P. Becher—contributing editor  w  Author to whom correspondence should be addressed. e-mail: fmonte@istec.cnr.it  Manuscript No. 22061. Received July 27, 2006; approved January 9, 2007.  Journal  J. Am. Ceram. Soc., 90 [4] 1130 - 1138 (2007)  DOI: 10.1111/j.1551-2916.2007.01589.x  r 2007 The American Ceramic Society  1130  \\x0c', 'The purpose of  this study was to measure and compare the  TSR of  two SiC-containing MB2-based ceramics, designed as possible base materials for structural  specifically  aerospace  TPSs.  In addition,  in-air oxidation tests were also performed.  Even though similar heat treatments do not fully reproduce the  real operative re-entry conditions, the results provided precious  complementary informations for the understanding of  the oxi dation behavior. The materials herein examined were already characterized in their radiative and surface catalytic behavior35:  thanks to low catalycity and high emissivity, they showed a good  potential for the selected application.  II.  Experimental Procedure  (1) Materials Processing  Two MB2-SiC mixtures of vol%,  commercial powders, amounts  in  ZrB2 þ 15SiC ðcomposition ZSMÞ ZrB2 þ 15SiC þ 10Hf B2 ðcomposition ZHSMÞ  were ball-milled in absolute ethyl alcohol for 24 h, using highpurity zirconia milling media. A quantity o5% of MoSi2 serving as sintering aid was batched into both the compositions.  Table I  shows  some characteristics of  the raw powders used.  After drying in a rotating evaporator, the powder mixtures were sieved through a mesh screen with 250 mm openings.  The powder mixtures were uniaxially hot-pressed in vacuum  using  an inductively heated graphite die,  lined with a BN sprayed graphitized sheet. Peak temperatures/dwell plied pressures were 18201C/15 min/30 MPa for ZSM, and 19001-19401C/45 min/40 MPa for ZHSM, about 201C/min av times/ap erage heating rate. The temperature was measured by means of  an optical pyrometer focused on the graphite die.  (2) Materials Characterization  (A)  Microstructure:  The bulk density (dB) and the theoretical density (dTH) were evaluated using the Archimedes method (water as immersing medium) and the rule-of-mixture,  respectively. The relative density (RD) was calculated dividing  the bulk density by the theoretical density . The phase compos ition was analyzed with an X-ray diffractometer (XRD, Ni-filCuKa  tered  radiation, model D500,  Siemens, Karlsruhe,  Germany) and a scanning electron microscope  (SEM, model  S360, Leica Cambridge, Cambridge, UK) combined with an en ergy-dispersive X-ray microanalyser (EDX, model INCA Ener gy 300, Oxford Instruments, Abington, UK). Polished sections  of the as-sintered materials were prepared with successively finer diamond-based abrasives ranging from 50 to 0.25 mm. A quali tative estimate of the grain size was made off from SEM images  of fracture surfaces.  (B)  Properties  (a)  Mechanical:  The Young’s modulus (E) and Poisson’s ratio (n) were measured on a 28 mm \\x02 8mm \\x02 0.8 mm plate using the resonance frequency method. Micro-hardness (HV1.0)  was evaluated by a Vickers indenter, using a 9.81 N applied load for 15 s. Flexure strength (s) in a 4-pt. configuration was tested  (model Z050, Zwick/Roell—Ulm, Germany) at room temperature (five specimens) on 25 mm \\x02 2.5 mm \\x02 2 mm chamfered bars using 20 mm and 10 mm as outer and inner span, respec tively, and a cross-head speed of 0.5 mm/min. The surface finish  of the bars, measured using a contact stylus type profiler (model  Talysurf Plus, Rank Taylor Hobson, Leicester, U.K.), was Ra 5 0.1470.01 mm, Ra being the arithmetic mean deviation of the assessed profiles. Fracture toughness (KIc) was measured ture (three specimens) using 25 mm \\x02 2 mm \\x02 2.5 mm bars on through the chevron notched beam method at room temperathe same jig used for the flexure strength (cross-head speed 0.02  mm/min). The bars were notched with an 0.08-mm diamond  saw;  the chevron notch tip depth and the average side length  were about 0.12 and 0.8 times the bar thickness, respectively. The ‘‘slice model’’ equation of Munz et al. 36 was used for the  calculation of KIc.  (b)  Thermo-Physical:  Thermal  expansion up to 13001C  was evaluated using a dilatometer (model DIL 402E, Netzsch Gera¨ tebau GmbH, Selb, Germany) in a stream of argon, 51C/ min heating rate. Thermal diffusivity (DTH) up to 12001C was measured through the laser flash method (model LFA-427,  Netzsch Gera¨ tebau GmbH), and following the  standard EN  821-2 as guideline. Heat capacity (CP) was measured using a modulated differential scanning calorimetry (model MDSC, TA  Instruments, New Castle, DE).  The TSR was studied through the method of the retained flexure strength after water-quenching (201C of bath temperature and 501C of thermal quench step). Chamfered bars 25 mm \\x02 2.5 mm \\x02 2 mm were used for testing. The experimental critical thermal shock DTC was determined using the guidelines outlined by the standard prEN820-3. In particular, DTC was identified using a linear interpolation between points that  first reduce the average flexure strength of the quenched bars by  more than 30% of the mean strength of the as-sintered material.  For  each quench temperature, at  least  three  specimens were  tested. In addition, the variation of the elastic modulus (EM) on the thermal shocked specimens was estimated using the load deflection method, and following the standard prEN 843-2 as  guideline. A cross-head displacement (recorded during the load ing cycle) from 50% of the peak load to 90% of the peak load  was selected in order to minimize strong non-linearities.  (c)  Resistance to Oxidation:  The resistance to oxidation was tested in flowing dry air (20 cm3/min) at 14501C for 1200 min, 301C/min of heating rate and free cooling, using a thermo gravimetric analyser (model STA449 Jupiter, Netzsch Gera¨ tebau GmbH, 10\\x003 mg of accuracy) equipped with a vertically heated Al2O3 chamber. Spacers of zirconia were placed between (dimensions 14 mm \\x02 2.5 mm \\x02 2 mm) and the the specimens Al2O3 holder with minimal contact area. The specimen’s mass was measured before and after exposure. XRD and SEM tech niques were used to analyze the microstructures of the oxidized  samples.  III.  Results  (1)  Microstructure Development  Table II shows processing parameters and some microstructure  features of  the hot-pressed materials. The bulk densities  dB of the as-sintered ZSM and ZHSM compacts were 5.61 and 6.06 g/cm\\x003, respectively. These values of density correspond to relative densities of 99% ZSM and 98.2% ZHSM. The fracture  surfaces observed by SEM displayed the  grain structure of  both the composites, with regularly faceted diboride (maximum size 5 mm). Representative regions of  grains  fracture sur Table I.  Characteristics of the Commercial Powders Used (Source: Companies Datasheet)  Formula  Company  Type  \\x003) Density (g cm  Particle size  Oxygen (wt%)  ZrB2 HfB2 SiC  H.C. Starck (Goslar, Germany)  B  6.09  o2 mm  2  Cerac Inc. (Bad Soden-Salmu¨ nster, Germany)  325 mesh  11.18  1.7 mm 11.6 m2/g 2.8 mm  w  —  H.C. Starck  BF 12  3.19  z  0.9  MoSi2  Aldrich (Milwaukee, WI)  6.26  y  1  w  Fisher size (APS).  z  Specific surface area.  y  Mean value.  April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1131  \\x0c', 'faces  revealed also a high percentage of  transgranular versus  intergranular fracture between the diboride grains in both the  compositions. The XRD analysis of the ZSM compact identified  the main phases ZrB2 and SiC, and an amount of monoclinic zirconia o1 vol%. Instead, in the ZHSM compact, obvious changes in the starting composition took place during sintering.  XRD analyses showed the formation of ZrB2/HfB2 solid solutions, hereafter indicated as (Zr,Hf)B2. Such solid solutions uphold the hexagonal structure of the initial ZrB2, but host at the same time Hf atoms. The presence of (Zr,Hf)B2 solid solutions, which very often appear as shells around ZrB2 cores, is exactly confirmed by the typical configuration of the diboride grains  structured  in  core-shell  (Fig.  1).  The  formation  of  such  (Zr,Hf)B2 s.s. was further assessed through localized EDX analyses on polished surfaces. Furthermore, the polished section  examined by SEM highlighted the SiC particulates distributed  intergranularly within the diboride matrices, in ZSM sometimes in agglomerates (maximum size 5 mm). At the same time, SEM  observations do not provide evidence of residual porosity.  Compared with the ZSM composition, the addition of HfB2 into the ZHSM composition enhanced refractoriness of the in itial powder mixture. In fact, higher peak temperature and long er hold were necessarily applied so far to complete densification: during the hold of 45 min, temperature drifted from 19001 to 19401C. Such conditions led the grain structure of  the ZHSM  composition to increase in the average size more than that of the  ZSM composition (Table  II). Moreover,  in ZHSM a partial  sintering  of  the  intergranular  SiC agglomerates  took  place.  Processing,  densification  behavior,  and microstructure  deve lopment  of  similar  systems  are  described  in  several  earlier  papers.9-11,21,24,37  (2)  Mechanical and Thermo-Physical Properties  Tables III and IV and Figs. 2-4 show the experimental data for  selected properties. Average values of flexure strength at room temperature (sR0) of 795 and 765 MPa for the as-sintered ZSM and ZHSM materials, respectively, result better than for a sim ilar system composed of ZrB2 mixed with 15 vol% SiC and then hot-pressed.37 The co-existence of very hard components like  MB2 and SiC, along with the fine and dense structure, favored high levels of micro-hardness. However, 4.1 MPa \\x01 m1/2 as mean value of fracture toughness in both the materials is typical for this family of (brittle) ceramics.2,9-12,24,37 The behavior of the retained strength (sR) after water-quench is shown in Fig. 2. The sR data of the thermal quenched bars in Fig. 2 indicate that a thermal shock severity DTC as high as 3851 and 4751C were sustained by ZSM and ZHSM, respectively,  before losing more than 30% of their initial mean flexure strength sR0 (Table III). Just for comparison, a ZrB2 material containing 15 vol% SiC (95% relative density and 235 MPa  mean flexure strength), tested using the water-quenching method, reported a TSR of 3501C.23 It should also be noted that the ZSM specimens quenched above DTC have an evident scatter of strength values above and below the reference threshold (i.e.,  70% of the initial mean strength). On the contrary, the ZHSM material seems to follow Hasselman’s theory,27 which predicts a  sharp drop in strength at a critical  thermal shock temperature  difference. The plots in Fig. 3 show the elastic modulus (EM) of a thermal shocked specimen against its sR. A substantial decrease in strength was very often accompanied with a significant re duction in elastic modulus.  In order to confirm that the decrease in flexure strength after  quenching was actually an effect of the thermal shock, four specimens of ZSM were heated up to 5001C (51C/min heating  rate), held at  that  temperature for 15 min to establish thermal  equilibrium, and then left cooling freely inside the furnace instead of quenching them into the water bath at 201C. The re850715 MPa mean71  tained  flexure  strength,  standard  deviation, measured at  room temperature on the heated (and  not-quenched) ZSM specimens, did not show significative difif compared with sR0 of III). This suggests that the decay in flexural strength measured  ferences,  the pristine material  (Table  after quenching is primarily caused by thermal  shock-induced  effects.  The conditions  for  failure in the pristine material during a  thermal quench is usually predicted by setting the maximum  permitted thermally-induced stress equal  to its tensile or bend  strength. Table IV lists  two TSR parameters calculated from  Eqs. (1) and (2)  R ¼ s ð1 \\x00 nÞ=ðaE Þ  (1)  R0 ¼ RKTH  (2)  where s, n, a, E, and KTH represent modulus of rupture, Poisson’s ratio, coefficient of thermal expansion, elastic modulus,  and thermal conductivity, respectively. The comparison of  the  TSR values in Table IV shows that the calculated R values are much lower than the experimental DTC.  Table II.  Processing Parameters and Microstructure Features  of the Hot-Pressed Materials: Peak Temperature (T ), Dwell Time (t), Maximum Applied Pressure (P), Theoretical (dTH) and Bulk Density (dB), Relative Density (RD), and Grain Size (GS)  Sample  Processing parameters  Microstructure features  T (1C)  t (min) P (MPa)  Density  GS (mm)  dTH (g/cm3)  dB (g/cm3) RD (%)  ZSM  1820  15  30  5.67  5.61  99.0  2  ZHSM 1940  45  40  6.17  6.06  98.2  3  Fig. 1.  Polished cross-sections of ZSM and ZHSM (SEM micrographs, secondary electrons); dark features are SiC particles. The core (C)- rim (R)  structure in ZHSM is indicated.  1132  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  \\x0c', 'Discrepancies between calculated R values and observed DTC have been reported previously by many investigators,25,26,33,34  and may indicate that cracks  initiate under  the so-called soft  thermal shock. These conditions imply that  the interior of  the  specimens cools before the surface, resulting in the need for a  stress reduction factor determined using the so-called Biot modulus b ¼ Lh=KTH (L characteristic specimen’s length, KTH thermal conductivity, and h surface heat transfer coefficient be tween specimen and cooling medium). The R parameter cap tures the initiation of thermal shock cracking under hard thermal shock conditions (i.e., values of b>2026): higher values  of R are in favor of greater resistance to fracture initiation during quenching. Differences between R and DTC are particularly evident for water-quenching thermal shock tests, primarily be cause of the dramatic reduction in heat transfer coefficient after  the formation of protective steam bubbles at the water/specimen  interface:  such a phenomenon would mitigate the severity of thermal shock, and result in overestimated DTC values, compared to the expected R parameter.33  The thermal conductivity (KTH) was calculated using the expression KTH ¼ DTHCP r, where DTH, CP, and r represent thermal diffusivity (Fig. 4), heat capacity (Fig. 4), and density, 301-12001C,  respectively. Within  the  interval  the  calculated  data indicate  that ZHSM tends  to dissipate heat  faster  than  ZSM (Table IV). Just for comparison, such values of KTH result lower than that reported for a ZrB2120%SiC, i.e., 103.8 W/ m1C at RT.38 In the present case of small sized specimens and significant KTH (i.e., low values of b), the TSR ranking is more adequately determined by the R0 parameter however, expected that in increasing L (i.e., the minimum heat  (Table IV).  It  is,  transfer  (specimen) dimension or its thickness in the case of the observed DTC values begin to approach the correspondent R parameters (see Table IV), mini standard flexure bar)  mizing the differences between the two tested compositions.  (3)  Resistance to Oxidation  The graphical trends of the specific mass change (w) versus time (t) during the 1200 min exposure at 14501C is plotted in Fig. 5, 3.4370.02 and 4.8870.02 mg/cm2 being the final w values for ZSM and ZHSM, respectively. Offsets equal to 0.4770.02 and 0.470.02 mg/cm2, which account for some oxidation prior the  1200 min hold, were respectively subtracted from raw data of  ZSM and ZHSM. Some data processing was done by calculating the dummy constant KD 5 w2/t points out that during the isothermal exposure obvious para (Fig. 6). Such an exercise  bolic kinetics (i.e., slope of the log KD equal to zero) were never established. In addition, the analysis of the very early stages of  the isothermal hold (see inset in Fig. 5) shows however prevail ing mass loss mechanisms over mass gain mechanisms, which in  turn become typical for the remaining exposure.  The XRD analyses on the surfaces exposed to air  revealed  monoclinic MO2, M 5 Zr for ZSM, and Zr/Hf for ZHSM, while the cross-sections emphasized a structure that consists of an  outermost glassy layer on top of a sub-scale composed of MO2 crystals (partially embedded within the same glass), which ex tends up to the unoxidized bulk (Fig. 7). The glassy layer, whose  composition analysed by SEM-EDX falls in the Si-O system,  adheres to the sub-scale, even though its thickness varies from  few to some tenths of microns. The values of about 50 and 100 mm are grossly representative of the oxide sub-scale thickness for  ZSM and ZHSM, respectively. In addition, EDX analyses as sessed that, in proximity of the interface between oxide sub-scale  and unoxidized bulk (Fig. 8),  sites  formerly occupied by SiC  particulates now appear occupied by a carbon-based solid com pound. Such an oxidation by-product, whose occurrence was already reported,10 was connected to the active oxidation of SiC,  which induces a partial depletion of SiC close to the inner ox idation front facing the virgin bulk. Such an oxidation mecha nism creates  so  far  porosity within  the  sub-scale  formerly  mentioned.  IV.  Discussion  (1)  Microstucture and TSR  Contrary to the well-known limitations connected to the densification of MB2-based materials,5,8-11,39 the obtainment of near full dense compacts was allowed by SiC which substantially en hanced densification of MB2 during hot-pressing. Similarly to powders of TiB2,40,41 MO2 and B2O3 were assumed as the main oxygen carriers upon the surfaces of MB2, M 5 Zr and Hf. Such a contamination by oxygen promotes vapor phase transport  (and thus coarsening) at temperatures below which mass trans fer mechanisms  like  boundary/volume  diffusion, which  are  much more effective than vapor phase for densification,  start  taking place:  the anticipated coarsening decreases  the driving  force for densification at higher temperatures. Densification of  SiC-containing MB2 powder mixtures peratures compared with pure MB2 are deemed to remove the oxide coatings separating MB2 particles from mutual contact.  initiates at  lower  tem 10,42 as  reactions with SiC  As  far as  the TSR evaluated through the retained strength superior DTC in the ZHSM material, compared with that of ZSM material, was at after water-quenching is  concerned,  the  tributed to more favorable heat transfer parameters conditions  Table III. Mechanical Properties of the Hot-Pressed ZSM and ZHSM Materials: Elastic Moduli E (Resonance Frequency Method) and EM0 (Load-Deflection Method), DE 5 (E-EM0)/E, Poisson’s Ratio n, Micro-Hardness HV1.0, Fracture Toughness KIc, and Flexure Strength at Room Temperature rR0  Sample  E  w  (GPa)  EM0  z (GPa)  DE (%)  n  zHV1.0 (GPa)  zKIc (MPaOm)  zsR0 (MPa)  ZSM  48074 50874  44475 47379  7.5  0.12  17.770.4 18.270.5  4.170.05 4.170.75  7957105 765775  ZHSM  6.9  0.128  w  Uncertainity.  z  Mean71 standard deviation.  Thermo-Physical Properties and TSR Parameters (R and R0 ) of the Hot-Pressed ZSM and ZHSM Materials: Linear Table IV. Coefficient of Thermal Expansion a (251-13001C), Thermal Conductivity KTH, and Critical Thermal Shock DTC  Sample  \\x006/1C) a (10  KTH(W/m \\x01 1C)  TSR parameters  301C  5001C  10001C  12001C  w  R (1C)  w  R0  (kW/m)  DTC (1C)  301C  5001C  ZSM  6.68  62.5  64.5  65.1  65.2  218  13.6  14.0  385  ZHSM  6.74  79.9  83.5  84.2  85.0  196  15.7  16.4  475  w  Calculated using sR0 in Eq. (1). TSR, resistance to thermal shock.  April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1133  \\x0c', '1134  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  Fig. 2. Retained flexure strength (sR) vs. thermal shock (T) of ZSM and ZHSM in water bath at 201C; full tested. Horizontal dashed segments: 70% sR0 of ZSM and ZSM and ZHSM (see Table III).  lines connect mean sR values at each T  that induce less severe thermal gradients across the specimen during the quench down test in a 201C water-bath. Being all the  tions, and are not solely dependent on materials characteristics.  Furthermore,  the determination of adequate  strength values,  specimens  tested of  the same (small) dimensions,  the different  such as the Weibull lower limiting strength,  is known to require  KTH (Table IV) was deemed to modify the actual Biot moduli of the two composites during quenching. It can at first be noticed to strength (s) ratios repre large fracture toughness (KIc) sent a favorable material’s characteristic for minimizing the ex that  tent of the crack propagation, but stands in direct contrast with  the requirement of high strain-to-failure necessary for preventing crack initiation.26,27,43 Hence, the combination of large E and s values which pertains the present materials may induce  deleterious damage resulting from fracture if it does occur upon severe thermal shocking, i.e., above DTC. It was in fact observed that, for thermal quench above DTC, a ZSM specimen spalled into a number of pieces, i.e., it had no measurable retained  strength. This situation very likely arose because of the sudden  release of  the large thermal strain-energy required to initiate a  crack from very small flaws in the material. The observed DTC for ZSM and ZHSM, respectively, seem to contradict the expected ranking based on the calculated R pa rameters. Lewis has  stressed the  importance of  selecting the  most appropriate value for strength so that meaningful rankings can be settled.34 In addition, Becher et al. argued that, as does  extensive testing. On the other hand, when extensive strength  data from appropriately sized specimens are unavoidable (this is  the present case), reasonable success can be obtained by taking a  characteristic strength as the mean value less two or three standard deviations.45 Calculated R values and experimental DTC under regimes of soft thermal quenchs have been repeatedly correlated through empirical expressions DTC 5 f(b) \\x01 R which try to approximate the thermal stress solutions for specific sample geometries of rod, plate, or bar.25,26,30,33,34 Also for the pres ent study, an empirical approach was adopted (damping factor f(b) 5 11B/b, 0o1/f(b)o1, b 5 Lh/KTH, h surface heat transfer coefficient, L minimum specimen dimension, KTH thermal conductivity, and B shape factor). According to the Lewis criterion,45 the mean strengths sR0 less three standard deviations (named sZSM and sZHSM in Fig. 9) were used for calculating corrected R parameters for ZSM (RZSM) and ZHSM (RZHSM), respectively. By setting L 5 2 mm, B 5 4, and KTH 5 83.4 W/ m \\x01 K at 4751C for ZHSM, an effective h of 67 kW/m2 \\x01 1C was inferred, being bD1.6 in correspondence of DTC equal to 4751C. This guess for the h parameter, though merely practical for such  the size dependence of the strength of ceramics, the temperature  a comparison,  falls anyhow very close to the typical  range of  dependence of thermal conductivity and surface heat transfer coefficient varies with temperature for different materials.33 Un fortunately,  surface heat  transfer  coefficients usually are not  known precisely and have been shown, moreover, to vary over  three orders of magnitude within the typical temperature range of water-quench tests. Therefore the actual DTC values for essentially all materials are function of the heat-transfer condi values for ceramics during water-quenching thermal shock tests,34,44 in particular 50 kW/m2 \\x01 1C for TiB2 at 5001C of thermal water-quench.44 Keeping the same values of h and L for the reduced KTH value of 64.2 W/m \\x01 1C at 3851C (that ZSM, means b approaching 2.1) induces more severe thermal gradients  in the quenched specimens, leading to a less effective capacity to withstand thermal stresses (i.e., DTC approaching 3851C). Being  Fig. 3. Retained flexure strength (sR) vs. residual elastic modulus (EM) of bars thermal shocked in water bath at 201C. Horizontal dashed segments: 70% sR0 of ZSM and ZHSM (see Table III).  \\x0c', 'April 2007  Resistance to Thermal Shock and to Oxidation of Metal Diborides-SiC Ceramics  1135  Fig. 4.  Heat capacity (CP) and thermal diffusivity (DTH) vs. temperature (T) of ZSM and ZHSM up to 12001C.  the dimensions of the specimens nominally identical, such a difference in DTC can thus be connected to the superior thermal conductivity that in the ZHSM material modulates a more favorable damping factor f(b) through the Biot number. It is also  scatter is associated to concurrent large scatter in retained elastic  moduli: there is no thorough explanation for this. On one hand,  to the  authors’ knowledge,  a  similar behavior  in analogous  systems was not verified. In ZSM material, the concomitant oc expected that, having both the composites levelled characteris currence of large scatter in retained values of strength and elastic  tics  like strength, Poisson’s  ratio, fracture toughness, thermal expansion and elastic modulus, the observed DTC parameters approach the calculated R values of about 2001-2201C in com ponents of larger dimensions.  In Fig. 3, the sharp reduction in elastic modulus EM was very often accompanied with a significant decrease in the retained strength sR. The general agreement between the reduction of flexure strength to coincide with a loss of elastic modulus makes  modulus  (Figs. 2 and 3) may be indicative of a non-uniform  growth in cracks density and/or  size due to sub-critical crack  propagation in the water quench bath. On the other hand, the  quench experiments have some well-recognized limitations. The  TSR tests herein adopted require for instance sample geometries  that are amenable to subsequent mechanical testing, and many  specimens should be tested to assure the significance of the re sults.  In addition,  the  specimens are tested destructively:  this  such combined measurements a useful  tool  to correlate the se implies that initiation sites and sub-critical thermal damage are  verity of the damage introduced during a thermal quench with  difficult  to identify. Thermal  stress gradients across  the small  the capacity to withstand further external loads. Experimental data in Fig. 3 clearly show that sR results more susceptible than EM to the influence of newly formed flaws during thermal shock. According to the Griffith relation among strength s, fracture toughness KIc and flaw size c, sEKIc = growth in size of just a single flaw after thermal shock have more obvious effects on strength than on elastic modulus.46 Evidence  the initiation and  p  ffiffiffi  c ,  of it is in both the compositions, when the elastic modulus declines less or not at all, compared with the corresponding sR values which exceeded an absolute loss of more than 30% sR0 (Table III). Therefore, for apparent drops of the elastic modu sized specimens may therefore contribute to the variability of the retained strength after quenching above DTC due to the probability of a initiation site being positioned in the high stress lo cations.  It  is known in fact  that  the flexure  strength testing  reaches the maximum stress only close to the specimen’s surface.  The combination of the thermo-mechanical characteristics in the  ZHSM material is somewhat more beneficial for a superior TSR  as compared with that of  the ZSM material, especially for the  small specimens herein tested. However, the absence of a large  scatter in retained strength and elastic modulus with increased  thermal shock seen for the ZSM composite offers a real advan lus, the corresponding strength decay is very likely connected to  tage for the ZHSM material. More extensive testing and char a substantial damage initiated and grown during thermal shock  acterizations would assist in identifying the difference in damage  in form of random network of cracks.  mechanisms and the benefits of such behavior.  A significantly larger scatter in retained strength differentiates  ZSM material from ZHSM material (Fig. 3). Also, the strength  (2)  The Resistance to Oxidation  The negative oxidation rate along the very early stages of exposure at 14501C (see inset in Fig. 5) discloses that the oxidation  of MB2  MB2 þ 5=2 O2 ðgÞ ¼ MO2 þ B2O3 ðl Þ  is however dominated by the volatilization of boron oxide  B2O3 ðl Þ ¼ B2O3 ðgÞ  (3)  (4)  from the  specimen’s  surfaces  exposed to the air-atmosphere.  Reactions 3 and 4 describe mechanisms of weight gain and  weight loss, respectively.  With increasing exposure time, the resistances to oxidation of  both the composites, due to SiC which is known to start oxi dizing more slowly than MB2, begin taking advantage from the ongoing formation of a glass coating on the faces exposed to air  (Fig. 7). Generally SiC, according to the reaction  Fig. 5. Specific mass change (w) vs. exposure time (t) during oxidation at 14501C. The inset expands the very early stages of the isothermal hold.  SiC þ 3=2 O2 ðgÞ ¼ SiO2 þ COðgÞ  (5)  \\x0c', '1136  Journal of the American Ceramic Society—Monteverde and Scatteia  Vol. 90, No. 4  Fig. 6.  Dummy parameter KD vs. exposure time (t), KD 5 w2/t.  reacts with oxygen and provides silica. Similarly to other ZrB2- SiC systems,1,3,6,10,13,15,16,24,47 significant amounts of silica com bine with the available B2O3 and yield a borosilicate glass. Such a glassy oxidation product is deemed to lose progressively the  B2O3 component, but is capable of providing better protection against oxidation than an MO2-based scale covered with only boron oxide.6 Although the borosilicate glass behaves as excellent oxidation barrier below 16001C,  it softens, due to the di minished viscosity. The undulating pattern of the external glass  is ascribed to the reduced viscosity.  The marked difference of  the thermo-gravimetric curves  in  Fig. 5 is still a matter of reasoning. The evaluation of the specific  weight changes could have been underestimated because of phe nomena connected to simultaneous gains and losses of mass,  which a previous work on the ZrB2-SiC system firmly asserted.15 The initial faster mass gain rate in ZHSM was attributed to  an inadequate provision of protective glass, compared to ZSM,  coupled with an easier diffusion of oxygen along short-circuited  paths through the forming (Zr,Hf)O2 sub-scale. Next, the reason why, compared to ZHSM, ZSM better resisted against oxida tion remains unclear. Having both the compositions the same  nominal  content of SiC and negligible levels of porosity,  the  core-shell structure of (Zr,Hf)B2 solid solutions in ZHSM supposedly was thought to offer less resistance to the inward dif fusion of oxygen, compared to ZrB2. Even though an oxidation process governed by the oxygen  diffusion through a glass that fits parabolic kinetics  seemed a  logic  expectation, departures  from an ideal parabolic pattern  were instead of not marginal significance (Fig. 6). A number of  factors is indicated responsible of such deviations: the evolution  of volatile products (B2O3, CO), non-steady state behavior due to faster oxidation of MB2 compared with SiC, and the active  Fig. 7. Fracture cross-section of the ZHSM material after 1200 min at 14501C in air  secondary electrons). An oxide sub (SEM micrograph,  scale (100 m thick) underlying the external glass (25 m thick) is indicated.  Fig. 8.  ZHSM material after 1200 min at 14501C in air; SEM micro graph, secondary electrons. SiC particles and carbon-based by products  (arrows) near the oxide sub-scale/unoxidized bulk interface are shown.  oxidation of SiC. The last factor remains being explained thor oughly. The examination through the SEM-EDX technique of  the  sub-oxide  scale/diboride  interface  (i.e  across  the  inner  oxidation front) revealed inclusions containing carbon (Fig. 8).  Being the size, shape, and distribution referable to the SiC par ticulates  in the unoxidized material,  these carbon-based struc tures were associated with an active oxidation mechanism of  SiC. Unlike some authors who claimed the complete volatiliza tion of SiC through active oxidation in SiC-containing MB2 matrices,6,15,24 the detection of C in such solid inclusions led to  consider the following thermodynamic equilibrium  SiC þ SiO2 , 2SiOðgÞ þ C  (6)  as the transition mechanism by which SiC starts experiencing the decomposition.48 The  SiC transforms  active  reason why  in  agreement with equilibrium 6 is not fully understood. Most re cently, Fahrenholtz developed a thermodynamic analysis of the ZrB2-SiC oxidation49: once SiC starts oxidizing according to reaction (7)  SiC þ O2 , SiOðgÞ þ COðgÞ  (7)  CO may further reduce to solid C at the so called sooting limit  (see reaction 8)  2COðgÞ , C þ O2 ðgÞ  (8)  Fig. 9. Biot number b vs. R parameter corrected via damping factor f(b) 5 114/b. According to Lewis criterion,45 sZSM and sZHSM, are sR0 of ZSM and ZHSM, respectively, less 3 standard deviations (see Table III). RZSM and RZHSM are obtained by using sZSM and sZHSM in Eq. (1).  \\x0c', 'In both cases,  it  is  as much as plausible  that  conditions  of high temperature and low oxygen partial pressure inside the the active oxidation of SiC.6,49,50  oxidized scale were met  for  Furthermore, once SiO(g) diffuses outward and encounters a  higher oxygen partial pressure, it would further convert into the  condensed silica phase. Glass pockets inside the oxide sub-scale  but  in the  vicinity of  the  external  glass may be  convincing  evidence of  the phenomenon just described.  It  should also be  considered that  the lack not only of open porosity but also of  easily-oxidizable secondary phases is a favorable condition that  does not offer preferential paths like pores or grain-boundary  channels to the inner transport of oxygen. As a consequence, the  oxygen partial pressure PO2 sufficiently low to induce the active oxidation of SiC. The PO2 parameter in fact is widely recognized as one of the primary  inside the bulk remains at a level  factors controlling the active-to-passive oxidation transition of SiC.48-51 Such an instability of SiC could have not negligible  influence on MB2-SiC systems in oxidizing conditions, dictating limited durability for prolonged services in thermally harsh ap plications.  V.  Summary  The present work explored the resistance to thermal shock and  to oxidation of two ultra-high-temperature MB2115 vol% SiC composites, M 5 Zr and Zr1Hf, brought to full density via hotrather uniform (2-3 mm  pressing. The microstructures were  grain size), with SiC incorporated intergranularly within the di boride matrices. The presence of SiC promoted the resistance to  oxidation of  the diboride matrices  through the coverage of a  silica-based glass which behaved as an effective barrier against  oxidation at high temperature. Both the compositions, oxidized at 14501C for 1200 min, had specific mass gains lower than 5 mg/cm2. Deviations  from parabolic  oxidation  kinetics were  seen, and attributed to evolution and release of volatile prod ucts, to non-steady state behavior based on the faster oxidation  of MB2 compared with SiC, and to active oxidation of SiC. The passive-to-active transition in the oxidation behavior of SiC  makes  the durability of  these SiC-containing diborides-based  composites in long-term thermally severe applications an issue.  The TSR,  tested through the  retained strength after water quench, showed that the ZrB2-SiC and (ZrB21HfB2)-SiC materials retained more than 70% of their initial mean strength (i.e., DTC) once the thermal shock has not exceeded 3851 and 4751C, respectively. On the contrary, the calculated R parameter the ZrB2-SiC and (ZrB21HfB2)-SiC materials were 2181 1961C, respectively. The improved DTC in the explained in terms of more  for  and  (ZrB21HfB2)-SiC was heat transfer parameters conditions  favorable  that establish less critical  thermal gradients across the specimens of reduced dimensions (i.e., bars 25 mm \\x02 2.5 mm \\x02 2 mm) during a quench down test. Having both the composites key characteristics like strength,  fracture toughness, elastic modulus, and thermal expansion very  similar,  the critical resistance to thermal shock of  larger speci mens is expected to approach the calculated R parameters.  Acknowledgments  The experimental activities were conducted within the frame of the Unmanned  Space Vehicle (USV) Italian National Program. The authors wish to acknowledge  the support of C. Melandri (ISTEC, thermal shock tests), A. Balbo (ISTEC, ox idation tests), G. Cosentino (CIRA, heat capacity tests), and M. Tului  (CSM,  Rome, Italy) for thermal diffusivity tests.  References  1K. Upadhya,  J.-M. Yang, and W. P. Hoffman,  ‘‘Materials  for Ultra-High  Temperature Structural Applications,’’ Am. Ceram. Soc. Bull., 58, 51-6 (1997). 2M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S. Johnson,  ‘‘Processing, Properties and Arc-Jet Oxidation of Hafnium Boride/Silicon Carbide  Ultra High Temperature Ceramics,’’ J. Mater. 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},{
  "_id": 237,
  "PDF": "Selective active oxidation in hafnium boride-silicon carbide composites above 2000°C.pdf",
  "Text": "['Journal of the European Ceramic Society 36 (2016) 3697-3707  Contents lists available at www.sciencedirect.com  Journal  of  the  European  Ceramic  Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Selective active oxidation in hafnium boride-silicon carbide composites above 2000  C  David L. Poerschke ∗ , Mark D. Novak 1 , Najeb Abdul-Jabbar, Stephan Krämer, Carlos G. Levi  Materials Department, University of California, Santa Barbara, CA 93106-5050, United States  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 22 January 2016 Received in revised form 23 May 2016 Accepted 29 May 2016 Available online 15 June 2016  Keywords:  Hafnium boride Hafnium carbide Silicon carbide Oxidation Ultra high temperature ceramic  The oxidation behavior of an ultra-high temperature ceramic (UHTC) based on HfB2 with 20vol% SiC and was studied following two 10 min arc-jet test cycles with nominal heat flux of 350 W cm−2 , stagnation pressure of 7 kPa, and a sustained peak surface temperature of 2360   C. Microstructure characterization 390  revealed a modified, layered structure comprising  \\u242em of porous HfO2 at the surface and an under740  lying  \\u242em porous region containing un-oxidized HfB2 over the bulk UHTC, unaffected below the oxidation front. The SiC presumably undergoes active oxidation, as commonly reported for temperatures above  1600   100   C. However, unlike typical of exposures below  2000   C no molten silicate phase was present at the surface to mediate the exchange of oxidant and gaseous reaction products. Additionally, a HfC impurity phase oxidizes concurrently with SiC rather than HfB2 . A thermodynamic analysis is provided to explain the observed behavior and the differences with lower temperature scenarios in the literature.  ±  © 2016 Elsevier Ltd. All rights reserved.  1.   Introduction  Ultra high temperature materials, defined as those capable of operating above 2000  C in aggressive environments, are enabling to critical technologies like hypersonic flight and rocket propulsion [1,2]. Progress  in these technologies  is thus  inextricably  linked to the capabilities of the materials, which  in many cases are not yet fully adequate in spite of decades of research and development. The menu of refractory compounds that have melting temperatures above 2000  C and stable oxides that also satisfy this criterion is presented  in Fig. 1. Preeminent examples are  the borides and carbides of Zr and Hf  [3]. While  the carbides have higher melting  temperatures,  lower coefficients of  thermal expansion  (CTE) [4,5] and are more elastically  isotropic owing  to  their cubic B1 structure, their hexagonal (C32) diboride counterparts have generally  received more attention because of  their higher  thermal conductivities  [6,7], and oxidation resistance  [6,8]. Nevertheless, the oxidation resistance of the pure diborides is inadequate above 1200  C owing to the volatilization of the nominally protective B2O3 amorphous layer produced during oxidation.  A favored approach is to add SiC, typically in contents of 10-30%, because the SiO2 formed during oxidation combines with the B2O3 to  form a more stable borosilicate glass The tradeoff  for the oxidation benefits with  the addition of SiC  is a marked decrease  in refractoriness. Although SiC  is  itself  reasonably  refractory,2 the attractive high melting point of HfB2 (TM = 3380  C) is substantially reduced by the formation of a eutectic at 2347  C even with small additions of SiC [10]. There are also eutectics at 3140  C between HfC (TM = 4215  C) and HfB2 [10], and at 2637  C between HfC and SiC [10], which suggests a possible ternary eutectic below 2347  C in the HfB2 -HfC-SiC system. [11]. It is also common to find impurity phases in the (Zr/Hf)B2 -SiC systems, notably (Zr/Hf)C in the boride powders [12], or WC incorporated during milling [13]. A  larger  fraction of  the research  in refractory boride systems has been performed on ZrB2 , presumably because of its lower density and  lower cost relative  to HfB2 . The  temperature capability, however, is significantly lower as the ZrB2 -SiC eutectic is 2207  C, a 1050  C  [10] and 140  C below  reduction over pure ZrB2 the HfB2 -SiC eutectic. In general, the oxidation behavior is qualitatively  ∗ Corresponding author. E-mail address: poerschke@engr.ucsb.edu (D.L. Poerschke). 1 Present address: ATI Wah Chang, Albany, Oregon.  http://dx.doi.org/10.1016/j.jeurceramsoc.2016.05.048 0955-2219/© 2016 Elsevier Ltd. All rights reserved.  2 Peritectic melting of SiC  is reported to form a Si-rich melt and solid carbon  in the range 2540   C to 2830   C while sublimation occurs at  2975   C at atmospheric pressure [9]. Sublimation may precede melting at lower pressures.                                              \\x0c', '3698   D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707  Fig. 1. Melting temperatures of refractory compounds versus that of their oxides limited to those materials wherein both exceed 2000   C.  in  included   in the ZrB2 systems will be   similar so  lessons  learned  the following description. Multiple approaches have been used to assess oxidation behavior of MB2 -SiC systems (M = Zr, Hf), namely furnace oxidation in air [2,14-16], resistive heating [17,18], fuel torches [19], arc-jet testing [12,20,21], and a scramjet wind tunnel [22,23]. At the lower temperatures (<1500  C) oxidation involves the formation of a surface layer of MO2 and a borosilicate glass [11]. When the layer is thin the CO resulting from the oxidation of SiC is assumed to be dissipated by dissolution into and diffusion through the glass. Concomitantly, continued oxidation is enabled by inward O2 diffusion through the glass  layer, and to some extent through the MO2 , which becomes oxygen deficient and  incorporates anion vacancies that  facilitate transport of the oxidant [24,25]. As the layer thickens and the oxidation  front moves  into the bulk of the material, the pressure of the gaseous oxidation products increases and occasionally bubbles burst through the glass, disrupting the oxidation barrier [26]. (Convection cells in the borosilicate phase have also been proposed as a mechanism that disrupts the outer scale [27].) At temperatures above 1500  C the glass layer stabilizes but the pressure of the gases evolved creates a porous  layer  free of glass between  the oxidation  front and  the  inner surface of  the glasspenetrated MO2, e.g.  [14]. At this point the SiC  is assumed to be oxidizing actively,  i.e. without the presence of a SiO2 layer at the surface of the particles, with the SiO(g) produced during oxidation condensing at the bottom of the glass covering the surface. A number of studies report this porous layer as containing the crystalline oxide of the metal on which the boride is based, e.g. [15,28]. There are, however, reports of the subsurface porous layer comprising primarily unoxidized MB2 , with only SiC being removed whereupon the  layer  is designated as a “depleted zone” [8,14,18,21].  In general, this behavior is reported in the literature for all experiments 2000  C, although a depleted zone is not always observed for up to  experiments performed at shorter times and/or lower temperature. The temperature regime above 2000  C  is much  less explored, largely because of experimental limitations. Arc jet testing has provided a path  to study phenomena  in  this extreme  temperature regime. In a notable set of experiments performed by Gasch et al. [12,20] it was found that under extreme conditions the glass layer is removed from the surface, leading to selective active oxidation of the SiC constituent, facilitated by a network of percolating channels that enable transport of the O2 reactant and the volatile products,  notably SiO, CO and B2O3 , with no hindrance by a condensed phase. The present investigation focuses on elucidating this regime of oxidation behavior, using specimens from the Gasch et al. study, kindly provided by NASA Ames for this purpose. Of particular  interest  is understanding the nature of the oxidation reactions at the respective SiC and HfB2 oxidation fronts and the conditions leading to SiC depletion without the formation of an outer SiO2 -based scale.  2. Approach  2.1. Material and testing  ±  ±  ±  ±  The material  investigated was an HfB2 -20vol%SiC UHTC  fabricated by NASA Ames as described  in detail by Gasch et al.  [20]. −325 mesh (<44  Briefly, the powders were  \\u242em) HfB2 (Cerac  Inc., \\u242em SiC (H.C. Starck, Newnow Materion, Milwaukee, WI) and 1-2  ton, MA). The powders were wet-milled with WC media, dried and hot pressed at 2200  C  for 1 h at 25 MPa, with  the chamber in <0.26 Pa vacuum up to 1600  C and then backfilled with 0.1 MPa Ar above 1600  C. The billet used to make the specimens used in this  0.01 kg m−3 , an elastic moduinvestigation had a density of 9.58  lus of 536   5 GPa, a Vickers hardness of 17   0.3 GPa, and a strength of 406   57 MPa [20]. The thermal conductivity of these materials 40 to  50 W m−1K−1 from ambient to 1800  C, and varied from  3.5 to 7 ppm K−1 from the coefficient of thermal expansion from  −100 to 1600  C [20]. Two specimens of the “flat face” type, 25 mm in diameter, were tested in the AMES Arc-Jet facility, both for 2 cycles of 10 min each [20]. Two conditions were applied, identified as “low” and “high”, involving calibrated heat fluxes of 285 and 350 W cm−2 and stagnation pressures of 5 and 7 kPa, respectively, with steady state surface temperatures of 1690  C and 2360  C, respectively, achieved during the second cycle  [20]. A SiC-coated graphite holder attached to a water-cooled sting arm supported the specimens  in the arcjet. This configuration  implies a  through-thickness  temperature gradient between the surface exposed to the arc  jet flow and the comparatively colder support. Because this paper deals only with the selective active oxidation leading to a SiC depleted zone, which is minimal  in  the “low” condition  [20],  the characterization and interpretation are focused only on the higher temperature specimen.  \\x0c', 'D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707   3699  2.2. Microstructure characterization  The specimen microstructure was characterized using a combination of optical microscopy, scanning electron microscopy (SEM; XL30 Sirion FEG, FEI, Hillsboro, OR) employing secondary (SE) and backscattered (BSE) electron imaging, and gallium ion imaging in a focused  ion beam microscope (FIB; Helios, FEI). The FIB was also employed  to extract  lamellae suitable  for  transmission electron microscopy (TEM) using FEI 200 kV Tecnai G2 Sphera and FEI Titan 300 kV FEG TEM/STEM systems. The  latter was used  for electron energy-loss spectroscopy  (EELS, Enfina 1000, Gatan, Pleasanton, CA) to determine the presence or absence of boron in select phases.  2.3.   Thermodynamic analysis  aO2 = 10−28  To aid  in  interpreting  the effects of spatial gradients  in oxidant potential and temperature on the microstructure evolution, the  thermodynamics of HfB2 , HfC, and SiC oxidation were considered. Calculations were performed using  the FactPS database in the Equilibrium module of the FactSage software package [29] (Thermfact/CRCT, Montreal, Canada). The oxidation of each  component was  simulated by equilibrating one mole of HfB2 , HfC, or SiC at 7 kPa over a  range of temperatures and oxygen activities (aO2 ) consistent with the “high” arc jet test conditions. The equilibrium phases and concentrations were  calculated  for all  relevant  species available  in  the FactPS database. This included 36 phases/species for SiC oxidation and 23 each for the oxidation of HfB2 and HfC. Argon was included in the calculations to maintain the gas pressure regardless of the presence of gaseous reaction products. Individual calculations at fixed 2400 values of  aO2 were repeated for  aO2 ranging from  aO2 = 10−4 , adequate to (sufficiently  low to preclude oxidation) to  oxidize the starting compounds without introducing excess unreacted oxygen  in the gas phase. These conditions are  intended to capture the effect of reduced oxidant activity within the gaseous boundary  layer and the evolving porous network rather than the free-stream conditions. While the quantity of B, C, Hf, and Si was fixed in each case, the number of moles of oxygen in the system was allowed to vary between successive  aO2 steps. However, except for a low O2 partial pressure in the gas phase to maintain the desired aO2 , the oxygen existed only in compounds with B, C, Hf, and/or Si. Note that while the calculations are described in the context of O2, at higher temperatures monatomic O  is the predominant oxygen species due to O2 dissociation. intervals of 10  C  to 25  C The calculations were  repeated at  between 1400  C and 2400  C to build a  library of phase equilibria data useful for the interpreting microstructure evolution within a temperature gradient and for comparison with previously published UHTC arc-jet [20,30] and  furnace tests [14,18,31] at  lower temperatures. The  total volume of  condensed  (solid or  liquid) phases was determined  for each calculation using  the  tabulated phase  fractions and the molar volumes of each phase relative to that of the oxidized constituent. Based on room temperature molar volumes, the HfO2 occupies 1.1 and 1.3 times the volume of the HfB2 and HfC consumed, respectively, and the molar volume ratios B2O3 :HfB2 , SiO2 :SiC, and C:SiC are 1.6, 2.1, and 0.4, respectively. This approach to examine the thermodynamics of oxidation and volatilization reactions  is complementary to the use of volatility diagrams for SiC [32] and ZrB2 [33] to understand the origin of the SiC depleted zone under a glass-permeated oxide  layer  in ZrB2 SiC UHTCs [31]. Specifically, the current analysis aims to capture the evolution of stable phases through the oxygen potential and temperature gradients between the  free-stream arc  jet flow and the various oxidation fronts present in the material.  Fig. 2. Microstructure of bulk HfB2 -SiC UHTC of the as fabricated microstructure, also seen away from the oxidation front (a) Back scattered electron image and (b) Bright field transmission electron micrograph. Small crystallites of HfC, lightest gray level in (a), are found dispersed within the main constituents, often associated with the SiC phase.  3. Results  3.1. Microstructure observations  The microstructure of  the bulk material was analyzed  in  the arc  jet  specimen, away  from  the oxidized  region, as well as  in small pristine specimens of the same ingot also provided by NASA. The typical features are shown in Fig. 2. In general, the specimens appear  fully dense, with no detectable porosity or microcracking  in  the areas examined. The BSE  image  in Fig. 2(a)  reveals the presence of polycrystalline aggregates of SiC, of order 10  \\u242em in size, with a small  fraction of HfC, an  impurity  formed during synthesis/processing [20]. The HfC grains are most commonly associated with the larger SiC regions. TEM analysis in Fig. 2(b) shows well-bonded  interfaces between all phases and  in some areas the presence of another minor phase tentatively identified as B4C (not shown). Fig. 3 shows an overview of the different zones  in the arc  jet tested with 2360  C steady-state surface  specimen  temperature. The entire cross section is shown in the inset, where the discolored rim represents the affected zone. The porous outer HfO2 layer was 390  \\u242em thick in the middle of the top surface, with some spallation in areas such as the upper corner shown in Fig. 3. The color of the outer oxide layer (not shown) changes from bright white near the surface to dark grey in the innermost portion. This color variation suggests a gradient in Hf oxidation state (i.e. concentration of oxygen vacancies in HfO2 in equilibrium with the local O2 potential) through the layer thickness, consistent with the known range of oxygen-deficient HfO2 stoichiometries at high temperature [24]. Between  the oxide and  the  remaining bulk material, which retained the microstructure shown in Fig. 2, there is a porous layer approximately 740  \\u242em thick containing only HfB2 , Fig. 4(a,c). There  \\x0c', '3700   D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707  Fig. 3. Cross section optical overview of the HfB2 -SiC UHTC button exposed to the arc-jet test with a nominal heat flux of 350W/cm2 for 2 cycles of 10 min each, with a steady state surface temperature of 2360   C. The button shown in the middle inset is 25 mm in diameter. The larger inset on the top right is an BS-SEM image of the corner, including both  the oxide  layer and  the depleted zone, superposed on  the  larger optical image that shows more clearly the different regions at this magnification.  was no evidence of SiC or a glassy silicate phase in either the oxide or the porous HfB2 layer, henceforth designated as the “depleted zone”. The oxide and depleted layers are also present on the sides of the specimen, albeit  in slightly  lower thicknesses, and even  in the backside, which was  in contact with the graphite holder—see Fig. 3. The thickness of these  layers along the backside decreases with distance away from the side of the button, suggesting gradual reduction in the oxidation kinetics, but even at the thinner end of these layers there was no evidence of a glassy phase in the pores of the oxide or the HfB2 . Closer views of  the depleted zone and  its  interfaces with  the oxide layer and bulk UHTC, denoted by dashed lines, are shown in Fig. 4. The scale of the porosity in the oxidized region is consistent with the size and distribution of the SiC phase in Fig. 2. Notably, HfC is still present in the remaining bulk material, Fig. 4(d), but not in the depleted zone, Fig. 4(c). The oxide and boride, which look very similar in contrast in Fig. 4(b), can be differentiated in part by the presence of extensive cracking  in the  former. Much  larger cracks normal to the surface are also periodically  found  in the depleted zone, Fig. 4(a), presumably associated with stresses induced during cooling. Fig.5(a)  reveals  that  the  removal of  the SiC occurs by active oxidation,  i.e. a SiC particle  is  receding without  formation of a condensed oxide,  leaving behind  a  void. The pore network  is sufficiently  interconnected  to  enable  continuous  ingress of O2 from  the  surface and  removal of  the SiO/CO gaseous products  Fig. 4. Cross section views of the modified UHTC microstructure after exposure to the arc-jet test. (a) Optical  image of the microstructure showing the SiC depleted zone between the remaining bulk material at the bottom and the oxidized layer at the top. The boundaries are defined by dashed  lines. The SEM images on the right correspond to the (b) oxide/depleted zone interface regions, (c) bulk of the depleted zone, and (c) the depleted zone/bulk interface regions. The darkest regions are voids, except near the bottom where they correspond to residual SiC, as noted in (d). There are small pockets of SiC completely surrounded by HfB2 which remain through the depleted zone.  by counter-diffusion. A minor  fraction of small SiC particles not connected  to  the  rest of  the percolating second phase network remains unoxidized as long as it is surrounded by the HfB2 matrix, as noted  in Fig. 4(d).  It  is also evident  in Fig. 5(b)  that  the HfC is oxidized very close  to  the advancing volatilization  front. The resulting microstructure  is highly porous and  too fine  to definitively determine the HfC oxidation sequence. However, preliminary microstructure and composition analysis suggests that the oxidation reaction may not involve direct conversion of the HfC to HfO2 and CO and instead proceed via an intermediate, carbon depleted HfC state, which would be consistent with previous observations at lower temperatures [34,35]. The transition between the HfB2 and HfO2 at the outer oxidation front exhibits some interesting features, as shown in Fig. 6. Gallium ion images in Fig. 6(a) and (b) reveal the boundaries between boride  \\x0c', 'D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707   3701  Table 1 Gas phase species observed at concentrations greater than 10−5 mol, presented from highest to lowest maximum concentration.  System   Species Type   Formula  SiC  HfB2  HfC  Oxide  Non-oxide  Oxide  Non-oxide  Oxide  Non-oxide   SiO, CO, CO2 , SiO2 Si, Si2 C, SiC2 , Si2 , Si3 , SiC, C, C3 BO, B2 O3 , (BO)2 , BO2 , B2 O B CO, CO2 Hf, C  The nature of the adherent and separated boride/oxide  interfaces is shown in detail in Fig. 6(c-e). In addition to the larger (of order 0.1  \\u242em) interface voids observed in Fig. 6(c) and Fig. 7, a population of smaller (or order 10(cid:5) s of nm) appears  in regions where the  interface  in nominally adherent, Fig. 6(e). EELS analysis along the path marked by a dashed  line across the  interface  in Fig. 6(e) reveals the abrupt transition in B content between HfB2 and HfO2 , Fig. 6(f). No B is detected in the HfO2 and there is not evidence of a transient B2O3 glass at the  interface. Additional analyses of the pores and thin cracks found at the interface showed no evidence of a glassy phase that might be associated with the presence of molten boria or a borosilicate glass. The bright field TEM imaging, Fig. 6(b), reveals heavily twinned oxide regions; diffraction analysis of the oxide  is consistent with monoclinic HfO2 . This structure suggests that the outer oxide layer grew as the high-temperature tetragonal phase and transformed to 1800  C [36]. This the monoclinic structure upon cooling below  conclusion  is further supported by the cracking within the oxide, presumably a result of the transformation strains.  3.2.   Thermodynamic analysis  Fig. 5. BSE  images near  the active oxidation  front at  the  interface between  the depleted zone and the bulk. Note the partially volatilized SiC particle and the evolving void space (V)  in (a), presumably percolating to the surface through the  fully empty pores within the HfB2 above. (b) Shows the partial oxidation of an HfC particle in a similar region.  and oxide much more effectively than the BSE image in Fig. 4(b). As expected, oxidation is proceeding from the boundaries of the pore network left by the earlier oxidation of the SiC. A layer of oxide surrounds the boride, which presumably recedes gradually until it is totally consumed. It is remarkable, however, that oxidation occurs over a relatively narrow band, of order 20  \\u242em thick, with no boride detected above it and no oxide below it in spite of the pore connectivity that provides access to the atmosphere. (See,  for example, that oxidation is proceeding from a pore on the right hand side of Fig. 6(b), single arrow, but not significantly from a pore immediately to its left at the same depth, double arrow.) Oxide may be found between boride particles, as shown both by the small oxide wedge penetrating through a boride grain boundary normal to the dashed  line  in Fig. 6(b), as well as between the HfB2 regions in TEM imaging in Fig. 6(c). To provide further insight into the oxidation process occurring at the HfB2 /HfO2 interface, a region near the advancing oxidation front was sectioned using the FIB. This series of 37 micrographs (with an average slice spacing of 0.2  \\u242em) captured the volume of approximately ten HfB2 particles surrounded by HfO2 . Sequences showing  the appearance of  two different HfB2 particles are shown in Fig. 7(a,b). The interfaces are partially attached, with thin high aspect ratio pores separating the boride and oxide around a portion of each HfB2 particle surface. As  illustrated  in Fig. 7, these thin  interface voids are observed to connect to the larger voids present in the HfO2 layer in at least one location for each HfB2 particle.  The thermodynamic calculations indicate that only a small fraction of  the dozens of  species  included  in  the  calculations  are stable at meaningful concentrations. To simplify  the  interpretation of  the  results, species occurring with mole  fractions below 10−5 10−7 )  (corresponding  to partial pressures below  for  all temperature/aO2 combinations  in a given system were excluded during post-processing. Reactions produce HfO2 , SiO2 , B2O3 , and C as condensed phases and the predominant gas phase species are summarized in Table 1. The evolution in stable phases and the volume of condensed phases for HfC, HfB2 , and SiC oxidation at 2360  C is plotted as a function of  aO2 in Fig. 8(b,d,f). For comparison, results of equivalent calculations at 1700  C are provided  in Fig. 8(a,c,e). These plots capture the general features of all calculations over the from 1400  C to 2400  C; the temperature-dependent pherange  nomena are further elaborated in the discussion. In the dynamic arc jet test environment, turnover of the atmosphere surrounding the specimen will replenish the oxidant and remove the gaseous reaction products allowing  for additional volatilization. These results therefore represent the stability of condensed phases  in the  limit of low gaseous flux. The oxidation sequence for HfC, Fig. 8(a,b),  is straightforward. aO2 HfC is the only stable phase. As the  At low  is increased the HfC oxidizes to a mixture of HfO2 and CO. This reaction results in a net increase in the volume of condensed phases according to the molar volumes reported in Section 2.3. With further increase in  the CO progressively oxidizes to CO2 . The primary change between 1700  C and 2360  C  is  the  required  to  initiate  the oxidation reaction. The oxidation of HfB2 ,  Fig. 8(c,d), generates HfO2 (with  an increase  in net condensed volume) and a combination of gaseous 1700  C  boron  oxide  species. At  the predominant  species  are B2O3 (g), (BO)2 (g), and BO(g). The latter two progressively oxidize to  aO2  aO2  aO2  \\x0c', '3702   D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707  Fig. 6. Details of the microstructure near the HfB2 oxidation front. (a, b) Gallium  ion  images show the transition from HfO2 near the top  into the remaining HfB2 in the depleted zone. Partially oxidized boride crystallites are embedded in the bottom layer of the oxide. There is an abrupt change between oxide growing from some pores at the interface but not from others immediately underneath. The BF TEM in (c) shows the boride oxide interface, which is separated in many regions. The twinned regions in the HfO2 provide evidence of the tetragonal to monoclinic transformation. TEM HAADF imaging in (d) and higher magnification in (e) show the substantially adhered interfacial region along the path marked by the dashed line in (b). EELS analysis along the scan marked in (e) is shown in (f), indicates no detectable B into the oxide.  B2O3 (g), At higher  aO2 a small fraction of B2O3 (l) contributes to the volume of condensed phases corresponding to the hump in the condensed volume plot, Fig. 8(c). Upon HfB2 oxidation at 2360  C BO(g) and  aO2 (BO)2 (g) initially predominate before oxidizing to B2O3 (g) and finally BO2 (g) at higher At 1700  C SiC oxidizes to a combination of CO(g) and SiO(g), Fig. 8(e). These reactions reduce the total condensed volume until the last SiC oxidizes, at which point there are no condensed phases. 10−13 )  If  the  is  increased  further  (beyond  the SiO oxidizes to  form SiO2 (cristobalite  is the thermodynamically stable phase at  this  temperature),  increasing  the volume of condensed oxide to more than twice that of the oxidized SiC. The CO subsequently oxidizes to CO2 . At 2360  C, the SiC oxidation sequence is more complex. Even at  aO2 sufficiently  low to preclude oxidation reactions, SiC sublimation forms a combination of SixCy species (Table 1) and graphite. The molar proportions of each species  in the gas phase are determined by the equilibrium partial pressures and the total pressure of the system. These species begin to oxidize as the  is increased. Carbon oxidizes preferentially leading to a further reduction in the fraction of SiC and an increase in the fraction of Si-rich gas species. These later oxidize to form SiO, which persists to high  aO2  aO2  aO2 .  4. Discussion  The microstructure  characterization  revealed  several  phenomena  relevant  to  understanding  the  performance  of  this temperatures above 2000  C. First, boride-carbide  composite at  active oxidation of SiC produces a thick SiC-depleted layer without forming an outer SiO2 -based scale, which would slow the oxidation kinetics. Second, the HfC oxidation front coincides with formation of the SiC depleted layer, suggesting a difference in the thermodynamic driving  force  for HfC and HfB2 oxidation. Finally, the HfB2 oxidation front features an abrupt transition between pores showing extensive boride oxidation and those without oxide growth. The processes giving rise to these microstructure  features are presumably influenced by gradients in temperature and  aO2 within the oxidation-affected zone. These gradients are expected to evolve as the (lower conductivity) oxidized layer grows and the respective oxidation fronts move deeper into the material (down the temperature gradient). The  thermodynamic calculations yield  information useful  for interpreting the role of these gradients in (i) the hierarchy of oxidation reactions with increasing oxygen activity and (ii) the tendency to  form condensed oxidation products contributing  to a protective oxide scale. The results of the thermodynamic calculations are  \\x0c', 'D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707   3703  aO2  ther increase in  aO2 , a portion of the SiO may condense as SiO2 . At 1525  C this transition occurs almost immetemperatures below  diately after oxidation while at higher temperatures  aO2 there is no condensed phase over an extended range of The key features from these maps are the (i) highest  aO2 at which a UHTC constituent  is thermodynamically stable and (ii) the  lowest  required  for  fractional SiO2 /B2O3 condensation at a given temperature. Curves corresponding to these transitions are plotted together  in Fig. 10.  It  is evident that at all temperatures HfB2 than HfC and SiC. Above 1650  C is more  resistant  to oxidation  aO2 than HfC while below 1650  C HfC oxiSiC oxidizes at  lower  dation is predicted to occur at  aO2 where some SiC remains stable. The results also indicate that condensed SiO2 can persist to higher temperatures than B2O3 .  4.3.   SiC depletion  Fig. 7. Sequences of images produced by FIB serial sectioning showing the distribution of micro voids (noted by black arrows) at the interface between HB2 particles and  the surrounding HfO2 near  the HfB2 oxidation  front. For each HfB2 particle observed, these thin  interfacial voids are  linked to the network of  larger voids  in at least one location, indicated here with white arrows. The slice spacing (into the plane) is indicated relative to the first slice.  first discussed in the broad context of understanding the oxidation process. Additional discussion focuses on the evolution of the specific microstructure features identified above in the context of both thermodynamic and kinetic considerations.  4.1.   Thermodynamic foundation  aO2  Based the calculated phase equilibria described in Section 3.2., the relevant  information  for  interpreting the oxidation and scale forming behavior for this multi-phase UHTC includes (i) the minimum  required to initiate HfB2 , HfC, and SiC oxidation, (ii) the tendency for SiC sublimation at elevated temperature, and (iii) the regimes  in which condensed SiO2 and B2O3 could contribute  to the formation of a protective scale.  In Fig. 18 these processes are represented by the curves  for the number of moles of HfB2 , HfC, HfO2 , SiC, SiO2 (cr,l), and B2O3 (l). To capture the processes over a range of temperatures and  aO2 , contours3 representing changes in the phase fractions were extracted from the relevant curves for each temperature. These contours were used to construct the maps in Fig. 9. Over the full range of temperatures from 1400  C to 2400  C HfC oxidizes to produce a mixture of HfO2 and gas, Fig. 9(a). The  required for HfC oxidation increases monotonically with temperature. Similarly, HfB2 initially oxidizes to form HfO2 and gas, Fig. 9(b). At temperatures up to 1700  C, further increase in the  leads to condensation of liquid B2O3 . Note that  nB2 O3 (l) is always less than 1 indicating the continued presence of boron in gas phase species. SiC sublimation is made evident by the decreasing SiC fraction aO2 (e.g., 10−27 ) in Fig. 9(c). The with increasing temperature at low  effect is most pronounced above 2100  C and by 2400  C less than 75% of the Si and C (on a gram atom basis) are present as SiC. As the aO2 is increased the SiC (or combination of and SixCy gases) oxidize aO2 for SiC oxidation shows considerably to CO and SiO. The critical  less temperature dependence than that for HfC and HfB2 . With fur aO2  aO2  3 Contours represent decadal changes in the phase content for subsets of the range between 10−12 and 0.99999 mol of the phase of interest. Additional contours added at 0.25, 0.5, 0.75, and 0.95 mol. For B2 O3 , a 0.98 contour is used instead of 0.99 for B2 O3 condensation since the latter falls outside the range captured by the maps.  aO2  As described  in the  introduction, much of the effort to understand the oxidation behavior of boride-carbide UHTCs has focused on temperatures below 2000  C with the majority of experiments performed below 1700  C [2,8,14,18,22,37]. Under these conditions simultaneous oxidation of  the boride and SiC produce an amorphous borosilicate  layer covering the surface. This scale acts as a barrier  to  further oxygen  ingress, depressing  the subsurface  and  facilitating the  formation of the SiC-depleted zone via active oxidation. At these lower temperatures the SiO oxidizes to SiO2 as it diffuses outward, up the  aO2 gradient. The condensation of SiO2 at the base of the borosilicate scale feeds the growth in scale thickness,  further  impeding oxygen  ingress and counteracting surface volatilization. This process is illustrated by the  ‘gap’ in condensed phases shown in Fig. 8(e). The width of the predicted gap without condensed SiO2 extends aO2 as the temperature increases. For temperatures below to higher  2000  C a significant  fraction of SiO  is expected  to condense as aO2 gradient. However,  SiO2 in  the outer region  the  the  fraction of condensed SiO2 would drop precipitously at higher  temperatures. The combination of temperature and  aO2 gradients in the arc jet experiment are such that the conditions needed to  form condensed SiO2 are not present within either the depleted zone or the outer HfO2 layer. SiO2 condensation below  the HfO2 layer, near the HfB2 oxidation front (i.e. at  aO2 below that for HfB2 oxidation), would require a temperature below 1600  C, Fig. 10. However, the observation that the oxide in the immediate vicinity of the boride shows evidence of the tetragonal   monoclinic phase transformation suggests  that  the  temperature at  the boride oxidation  front should have been above 1830  C [24]. SiO2 condensation within or above  aO2 the HfO2 layer could occur at higher temperatures, but would require correspondingly higher The high vapor pressures of non-oxide  species above SiC at 2360  C,  i.e. the sublimation noted  in Fig. 8(f) and Fig. 9(c),  likely further  facilitates  the  formation of  the SiC depleted zone during the early stages of the arc jet exposure.  →  4.4. HfC oxidation  At high temperatures (e.g. 2360  C) HfC  is expected to oxidize at an oxygen potential between that  for HfB2 and SiC oxidation, Fig. 10. Based on the presumed  aO2 gradient through the depleted zone, this sequence would produce an outer depleted  layer containing only HfB2 and an inner depleted layer containing both HfB2 and HfC.  Instead, the advancing HfC oxidation and SiC depletion fronts are coincident  in the “high” condition of the arc-jet exposure, suggesting that the processes are  linked. Given that the HfC is usually found in the vicinity of the SiC, the removal of the latter by active oxidation exposes the HfC to the atmosphere within the porous network, allowing HfC oxidation to proceed, e.g. Fig. 5(b).  \\x0c', '3704   D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707  Fig. 8. The condensed volume corresponds to the total volume of condensed phases at a given condition (computed using molar volumes and the phase fractions at left), normalized by the volume of one mole of the species of interest.  Based on the stability of the surrounding HfB2 It is apparent that HfC is oxidizing at a lower  aO2 than HfB2 , consistent with the hierarchy predicted by the thermodynamic calculations, Fig. 10. Two potential explanations for the observed microstructure are considered. The first scenario  is that the active oxidation of SiC  is limited by the egress of the reaction products rather than oxygen ingress, with a concomitant rise in the local  aO2 to a level sufficient for simultaneous SiC and HfC oxidation. Alternatively,  the  local conditions dictated by the through-thickness temperature gradient could allow that HfC oxidation to proceed at the same (or  lower) aO2 than SiC. These scenarios are discussed below. The UHTC oxidation kinetics has been considered for the refractory boride alone [38] and composites of the boride and SiC [39]. While these models effectively predict the behavior at lower temperatures, they did not explicitly consider the behavior ensuring a complete loss of the glassy SiO2 -based layer. In the absence of a dense scale, the kinetics  is presumably dictated by gaseous diffusion through the porous HfO2 and HfB2 layers. For the present case  molecular diffusion is expected to be dominant, based on standard criteria [40] and assuming (i) a mean temperature of 2000  C, (ii) a gas comprising predominantly N2, and  (iii) a pore neck diameter of order 0.5  \\u242em (selected for consistency with previous work [38,39]). Given the nearly equivalent O2 and CO diffusivities4 under these conditions  [41],  the  relative fluxes are determined by  the concentration gradients driving diffusion. The oxygen concentration gradient is constant, fixed by the free stream concentration and the low oxygen potential at the advancing oxidation front. The magnitudes of the CO and SiO gradients are determined by the equilibrium constants for the oxidation reaction at a specific  Under  conditions where  SiC  oxidation  is  unfavorable—i.e. 10−20 , Fig. 9(c)—the equilibrium CO and SiO concentrations will remain  low,  limiting the driving  force  for outward diffusion,  aO2 <   aO2 .  4 The SiO diffusivity is likely lower owing to its larger size and mass, but presumably within a factor or two.  \\x0c', 'D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707   3705  Fig. 10. Comparison of  temperature-dependent  tendency  for oxidation of UHTC constituents and formation of condensed oxides, based on data presented in Fig. 9. Curves for oxidation represent the highest  aO2 at which HfB2 , HfC, or SiC is stable. Curves for condensation correspond to the first appearance of SiO2 and B2 O3 with increasing   aO2 .  (Fig. 10) the thermodynamic driving force for HfC oxidation would be greater  than  that  for SiC oxidation; even  lower  temperatures could establish conditions where oxidation product egress was \\x01T of at least 700  C between rate-limiting. These scenarios imply a  the outer surface and the SiC oxidation front. In the quasi-steady state condition the top surface temperatures is 2360  C and the HfB2 oxidation front must be above 1830  C based on  transformation observed  in  the HfO2 . Given  the oxide  thick\\u242em and nominal heat flux of 350 W cm−2 , the estimated ness of 390  effective thermal conductivity of the porous HfO2 layer would be 2.6 W m−1 K−1 . This result  is reasonable  in the context of thermal conductivity data for monoclinic ZrO2 and  is about twice the value estimated  for  the minimum  thermal conductivity of HfO2 from existing models [42,43]. The temperature drop through the depleted layer should scale with its thermal conductance relative to that of the HfO2 layer. The thermal conductivity for pure, dense is of order 70 W m−1 K−1 20% porosity HfB2 [3]. Corrected  for  (assumed spherical) resulting from SiC volatilization, the depleted 50 W m−1 K−1 . With a thickness of layer conductivity would be  740  \\u242em the thermal conductance of the depleted layer would be about 10 times that of the oxide  layer, giving a temperature drop 50  C. Combined with the  through the  former of  \\x01T across the oxide, the estimated temperature of the SiC active oxidation front is of order 600  C5 below that of the surface. This value is slightly less 700  C  than the  \\x01T derived from Fig. 10, but, given the approximations involved, it is possible that the local conditions could support simultaneous SiC and HfC oxidation. This scenario becomes more plausible if one considers that HfC oxidation sequence does not  involve direct conversion of HfC to HfO2 . Possible scenarios include dissolution of oxygen into the carbide (well documented  in ZrC [44]), depletion of carbon from the 65 at% Hf [45]), an alternate carbide (the carbide  is stable up to  intermediate  reaction product  (observed at  lower  temperatures [35]), or  initially oxygen-deficient HfO2 . These  intermediate pro 5 Due to the large disparity between the relative conductivities, this result is not particularly sensitive to either absolute conductivity value.  Fig. 9. Maps summarizing the temperature and  aO2 depence on the stability of (a) HfC, (b) HfB2 , and (c) SiC and condensed (b) B2 O3 and (c) SiO2 . Description of the salient features is provided in Section 4.  aO2 >   and creating conditions where oxidation-product egress could be rate controlling. However, conditions leading to extensive SiC oxi10−16 ) would generate CO and SiO gradients dation  (i.e.  similar  in magnitude  to  the  inward oxygen gradient. As such,  it is expected that the oxidation kinetics could only be controlled by 10−16 . the outward flux in the case of  aO2 significantly below  aO2 for HfC oxidation decreases According to Fig. 10, the critical  with temperature more rapidly than that for SiC oxidation.  1650  C If the temperature at  the SiC depletion  front were  less  than   \\x0c', '3706   D.L. Poerschke et al. / Journal of the European Ceramic Society 36 (2016) 3697-3707  aO2  cesses are  likely  to occur at  lower  than  those calculated  for the direct conversion to stoichiometric HfO2 . Due to the fine scale of the observed microstructure, the present characterization was unable to conclusively determine the HfC oxidation sequence but the evidence suggests that HfO2 is not the initial reaction product. Nonetheless, the collective evidence and analysis suggest that the combination of the  imposed thermal gradient and nuances of the HfC oxidation process allow for concurrent SiC and HfC oxidation when the latter is exposed to the gas phase by the volatilization of the former.  4.7. HfB2 oxidation  In the (ultra-high) temperature range of interest for this study HfB2 oxidation should yield solid HfO2 and gaseous BxOy species, Figs. 8 and 9. Continued HfB2 oxidation requires both inward oxidant transport and the egress of these gaseous oxidation products. The porous structure of the HfO2 scale presumably facilitates transport through the bulk of this oxide  layer. At the boride oxidation front, Fig. 6(a,b), the oxidation process appears to proceed via HfO2 growth from the pore walls, encircling HfB2 particles with oxide. However, the oxidation of HfB2 at each successive level appears to proceed at a rate similar to the inward progression of the oxidation front.  In other words, the transport between the oxide-encircled boride and atmosphere proceeds with similar kinetics to  inward oxidant transport down the  aO2 gradient. Some oxygen transport could occur through the HfO2 in the solid state, particularly given the high vacancy content  implied by the darker coloration of the innermost HfO2 . However,  the kinetics cannot be explained by solid-state diffusion alone, particularly  in  light of the negligible B solubility in the HfO2 . Instead, it is likely that the thin, surface connected voids at the boride/oxide interface form and are maintained to facilitate BxOy egress.  4.8.   Implications for UHTC application  Many of the proposed applications for UHTCs require service at temperatures above 2000  C where, based on the analysis herein, a condensed SiO2 -based outer scale is unlikely to be stable. As such, the benefit of improved oxidation resistance afforded at lower temperatures via SiC additions to refractory borides are unlikely to be realized at higher temperatures. Instead, the SiC is rapidly depleted from the composite, arguably reducing the mechanical integrity of the remaining SiC-depleted boride  layer. Additionally,  the  (relatively) low temperature eutectic between SiC and other refractory borides and carbides significantly reduces the maximum use temperature of multi-phase UHTC systems employing SiC as a silicon source. Alternate materials strategies involving modified additive combinations are needed to increase the stability of a SiO2 -based scale [46] and/or  reduce  the  selective oxidation/depletion of  specific phases (i.e. SiC) at high temperatures. Preliminary work on systems containing transition metal silicides  instead of SiC show promise [47,48] but additional work is clearly needed in this area.  5. Conclusions  The present investigation has elucidated the transition between two different scenarios for the selective active oxidation of SiC  in HfB2 -SiC UHTCs  leading  to  the  formation of a SiC depleted zone under a porous outer HfO2 layer. In the more commonly observed scenario  there  is a continuous  layer of  (boro)silicate glass filling the pore structure of the outer layers that regulates the ingress of oxidant and the release of the gaseous reaction products. In the second scenario, emphasized in this work, the combination of higher temperature and oxygen activity precludes  the condensation of  the silicate  layer, enabling the direct counterdiffusion of oxidant and reaction products through the network of pores created by the selective oxidation of the percolating SiC network. It  is  inferred  from  the microstructural  evidence  that high temperature oxidation  in  the arc-jet proceeds with  two distinct reaction fronts moving  into the UHTC. At the faster moving front the oxygen activity and temperature enable the concurrent oxidation of HfC and the direct conversion of SiC  into gaseous species but are below the threshold to oxidize the HfB2 matrix. The second front, closer to the surface, has a temperature and oxygen activity sufficient to oxidize the porous HfB2 but not to enable the condensation of SiO2 , negating the  formation of glass within the porous (9%) associated network. Except  for  the small volume  increase  with the conversion of HfB2 into HfO2 , and perhaps some sintering, the pore network left by the volatilization of SiC is largely retained during oxidation, providing paths  for gaseous diffusion between the active oxidation front and the surface.  It  is especially notable that the oxidation of HfB2 occurs within a very narrow band, with no boride observed above  it and no oxide below  it. The oxidation of the boride involves transport of O2 and removal of gaseous B2O3 through a network of thin pores evolving at the interface between each HfB2 particle and the surrounding oxide, which are  in turn connected to the larger pore network of the volatilized SiC. Thermodynamic calculations provide insight into the effects of temperature and oxygen activity on the oxidation of the  individual UHTC components, confirming the scenarios above. Maps were developed to define the regimes of stability of different condensed and volatile species, the conditions needed for concurrent oxidation of HfC and SiC, and those leading to the formation or absence of a silicate layer over the depleted zone. The analysis is generally consistent with prior reports in the literature for similar phenomena.  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},{
  "_id": 238,
  "PDF": "Separating Test Artifacts from Material Behavior in the Oxidation Studies of HfB2-SiC at 2000 degrees C and Above.pdf",
  "Text": "['Int. J. Appl. Ceram. Technol., 10 [2] 293-300 (2013)  DOI:10.1111/j.1744-7402.2011.02730.x  Separating Test Artifacts from Material Behavior in the -SiC at 2000°C and Above Oxidation Studies of HfB2  Carmen M. Carney* and Triplicane A. Parthasarathy  UES, Inc, Dayton, Ohio 45432  Michael K. Cinibulk  Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright Patterson Air Force  Base, Ohio 45433  Oxidation characteristics of HfB2-15 vol% SiC prepared by field-assisted sintering was examined at 2000°C by heating  it  in a zirconia-resistance furnace and by direct electrical  resistance heating of  the sample. Limitations of  the material and the  direct electrical resistance heating apparatus were explored by heating samples multiple times and to temperatures 2300°C. Oxide scales that developed at 2000°C from both methods were similar  in that  they consisted of a SiO2/HfO2 outer  in excess of  layer, a porous HfO2 layer, and a HfB2 layer depleted of SiC. But  they differed in scale thicknesses,  impurities present,  scale  morphology/complexity. Possible test artifacts are discussed.  Introduction  resistance compared with the diborides alone near 1600° C.1-3 At temperatures below 1800-2000°C, a refractory  Characterization of  the oxide scale formed on ultra  porous metal oxide scale is formed that  is protected by a  high temperature ceramics (UHTCs) has been a topic of  glassy silica scale. However, as temperatures are increased  intense study over the past decade. In particular, compos the protective  silica becomes  less  viscous  and thus  less  ite  systems  of  diborides  of  hafnium or  zirconium  protective. Hypersonic flight will  require  leading edges  with SiC have been studied for their improved oxidation  and nose  cone  components  and cooling to temperatures  to withstand rapid heating in excess of 2000°C under  This  study was  supported by the United States Air Force Contract # FA8650-10-D-5226.  *carmen.carney@wpafb.af.mil © 2013 The American Ceramic Society  shear stresses imparted by air flow. In addition, the envi ronment within the boundary layer near the component  \\x0c', 'will be comprised of a fraction of dissociated atoms depending on velocity, altitude, and other factors.4,5 Dis sociated oxygen can alter  the boundary between active  and passive oxidation of SiC and thus kinetics.6 These  influence oxida tion  temperatures  and  conditions  are  unattainable  in traditional molybdenum disilicide  ele ment  furnaces whereas  conventionally  accepted arc  jet  testing is expensive and the primary oxidant is dissociated  oxygen, so new methods of testing have been developed. Methods such as oxyacetylene torch heating,7-9 laser heatthe sample itself,11-13 ing,10 direct resistance heating of and scramjet simulators14 are being developed. The rapid  heating profiles and higher  temperatures attainable with  these tests may lead to different oxide morphologies and  performance  than  those  observed with  furnace-heated  samples. It is imperative that a correlation between differ ent  testing methods is made so that samples prepared by  different  exposure methods may be compared. To this -SiC samples were heated in air at 2000°C  end, HfB2  using a  zirconia-resistance  furnace  and direct  resistance  heating of  the sample and the resulting oxidation prod ucts were compared. In addition,  the limits of resistance  heating including multiple cycles and maximum temper atures were examined.  Experimental Procedure  Commercially available HfB2 (Materion, Milwaukee, WI, 99.9%, 45 lm) and b-SiC (Materion, 99.9%, 1 lm) were used to prepare HfB2-15 vol% SiC (HS).  The powder mixtures were ball milled in isopropanol  for 24 h with SiC grinding media, dried at  room tem perature, and subsequently dry milled for 12 h. Typical  weight  loss of  the SiC grinding media after milling was  0.2 mg (0.2 wt% of the total batch). The were sieved through an 80-mesh (177 lm)  powders  screen.  A quantity  of  150 g  of  the milled  powders was  loaded  into  a  60-mm diam.  graphite  die. A layer  of  BN and graphite  foil  separated the  powder  from the  die with the powder  in contact with the graphite  foil.  The  powder-filled  dies were  cold  pressed  at  approxi mately  50 MPa.  The  powders  were  sintered  using  field-assisted  sintering  (FAS: HPD 25-1, FCT 2000°C for  Sys teme, Rauenstein, Germany)  at  15 min  under  a  32 MPa  load. The controlled rates were 50°C/min. The load was applied to 1600°C and released on cooling  heating  and  cooling  during heating 1000°C.  to  Oxidation samples were cut with a wire electro-disinto 5 mm 9 5 mm 9 3 mm and 53 mm 9 3.5 mm 9 19-25 mm long  charge machine (EDM)  rectangles  (furnace heating)  5.0 mm rectangles with  a  centered  3 mm thick region of  reduced area (resistance heating).  The samples were polished using diamond slurry 1 lm finish.  to a  Polished  samples were  heated  by  a  zirconia-resis tance  furnace  (ZrF-25: Shinagawa Refractories, Tokyo,  Japan)  and direct  electrical  resistance. Macrographs of  the two tests and sample geometry are shown in Fig. 1.  The  furnace heating was  accomplished by a molybdeto 1100°C and a zirconia ele5°C/min. The  num disilicide pre-heater 2000°C at  ment  to  a  rate  of  samples  were  held  at  temperature  for  30 min.  Temperature  measurements  were  performed  using  a  single  color  pyrometer  focused  on  the  zirconia  element.  Samples  were  supported  on  a  zirconia  crucible. The  zirconia  crucibles  were  cut  from a  larger  crucible  (Advalue,  Tuscon, AZ; 10 mL Ca-stabilized ZrO2 ZrO2 and 4 ± 1% Ca). Direct electrical  crucible; 95%  resistance heat ing was  controlled  by  the  power  output  of  an AC  power  supply  across  the  sample  and  temperature was  read by a two-color pyrometer  (FMP2; FAR Associates,  Macedonia, OH)  that was  focused on the center of  the  reduced-thickness area. The samples were held in place  between two graphite  spacers by  tightening  set  screws  on  the  copper  electrodes. Ag  paint was  used  on  the  ends of  the samples  to improve electrical contact. Tem perature, current, and voltage data were recorded using  LabVIEW (National  Instruments, Austin, TX). Table I  lists  the oxidized samples with their heating conditions.  Oxidized samples were mounted in epoxy and pol ished in cross  section perpendicular  to the bottom (side  facing the crucible or notched side) of the sample to a 1 lm finish using diamond slurry. The microstructures  were  characterized  using  scanning  electron microcopy  (SEM, Quanta,  FEI, Hillsborough, OR)  along with  energy  dispersive  spectroscopy  (EDS,  Pegasus  4000;  EDAX, Mahwah, NJ)  for  elemental  analysis. All EDS  analysis was done using 15 kV accelerating voltage and  at  least a 100 s  live capture time.  Results  Single 2000°C Exposure  The heating profiles of  the HfB2-15 vol% SiC zir conia-resistance  furnace  heated  sample  (HS-F)  and  294  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  \\x0c', 'direct  electrical  resistance-heated  (HS-R)  sample  are  shown  in Fig. 2a  and  b, respectively. The maximum the HS-R sample was 2027°C  observed temperature of  using 82.5 V and 20.3 A (averaged over  the hold). The  oxidized  HS-F  sample  had  a  thicker  oxide  scale  (Fig. 3a) compared to the HS-R sample (Fig. 3b). The  HS-F sample was exposed to oxidizing temperatures  for  a greater length of time than the HS-R sample (6.5 h 1100°C compared ~4 min 800°C).  above  to  above  The oxide layers  labeled in Fig. 3a (HS-F) and Fig. 3b  (HS-R)  are  composed  of  (I)  a  SiO2-based  glass  that  penetrates  a HfO2-based skeleton;  (II)  a porous HfO2  scale;  and (III)  a SiC-depleted layer. The SiC-depleted  layer  is  defined  as HfB2 with a  reduced SiC content  (partially  oxidized SiC). The  average  total  oxide  scale  thickness measured from the top side of the HS-F sample is 660 ± 45 lm with the depleted layer comprising  53% of  the  scale. The  thickest  total oxide scale mea105 lm with 5% of  sured on the HS-R sample was  the total  scale consisting of  the depleted layer.  The oxide scales of HS-F samples possess a distinct  two-phase SiO2-based glass with the  less-pure  (less vis cous)  impurity-laden glass  rising  to the  surface of  the  oxide scale and the purer glass  found deeper within the  scale  (Fig. 4a).  A  two-phase  glass  found  in  furnace  heating has been described previously by  the  authors,  which was rities.15  shown to contain Al and Ca as major  impu In addition, HfSiO4  (with  a Ca  impurity)  is  found in the HS-F sample, but not  the HS-R sample.  The existence and absence of HfSiO4 was confirmed by  XRD. Figure 4b  is  an EDS  comparison of  the purer  (darker) and impure  (lighter) glasses  in the HS-F sam ple  along with the HfSiO4  and HfO2 phases.  In the  HS-R sample Al  impurities can be found randomly dis tributed  throughout  the  glassy  phase  (inset  Fig. 4c).  Figure 4d is  a  representative EDS spectra  of different  locations within the HS-R glass. There  is no hierarchy  to  the  concentration  of Al  in  the  glass  phase when  comparing  the  chemistry of  the glass  along the  length  of  the HS-R oxide scale.  Repeated 2000°C Exposure  An  advantage  of  the  direct  electrical  resistance  heating test  is  that  the sample can be exposed multiple  times  to the  same or different heating profiles. A sam2000°C twice  ple  (HS-Rr) was  heated  to  using  the  same heating profile as HS-R. The maximum observed 2030°C. The  temperature was  heating  profile  and  a  (a)  (c)  (b)  Fig. 1.  (a) Macrograph of direct electrical  resistance heating  apparatus. The temperature is read by means of a fiber optic  cable through a carbon tube (1)  that  leads  to the pyrometer.  Power is  supplied through the copper electrodes  (2)  to a heated  sample (3) gripped by carbon spacers  (4).  (b) Macrograph of  the  zirconia-resistance furnace showing the sample stand (A),  cylin drical zirconia heating element  (B), alumina insulation (C), and  molybdenum oxide insulation (D).  (c) Samples prepared for  direct electrical heating (top) and zirconia-resistance heating  (bottom) on the supporting zirconia crucible. The Ag paste  improves electrical  conduction.  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  295  \\x0c', '296  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  Table I.  List of  the Oxidized HfB2-15 vol% SiC Samples and their Heating Conditions  Sample ID  Test method  Max. observed temp. (°C)  Hold time (min)  Comments  HS-F  HS-R  HS-Rr  HS-R2  HS-Rf  Furnace  Self-heating  Self-heating  Self-heating  Self-heating  2000  2027  2030  2041  2325  30  1  1  2  0  —  —  Two 1 min holds  —  Heated to failure  Fig. 2.  Heating profiles of  the (a) HS-F (calculated) and (b)  HS-R (actual)  samples.  micrograph  of  the  resulting  oxide  scale  are  shown in  Figs. 5a and b. The oxide scale has a periodic structure  consisting  of  layers  of  SiO2  and HfO2 penetrated by  SiO2. For comparison, 2000°C  2 min.  for  to  observed  temperature  of  sured oxide scale of  a  sample  (HS-R2) was heated  (Fig. 5c) with 2040°C. The  a maximum  thickest mea found for  the HS-R sample  the HS-R2 sample was double that (217 vs 105 lm), and the  oxide  scale was not  composed of periodic  layers. The  lines.  the HS-F sample after oxidation at  Fig. 3. (a) Micrograph of 2000°C (b) Micrograph of the HS-R sample after oxidation at ~2000 °C. The oxide layers are (I) HfO2 penetrated by SiO2, (II) porous HfO2, and (III) depleted HfB2 layer. The approximate boundary between layers is shown by the dashed white  \\x0c', 'oxide  scale  formed near  the  center of  the  reduced area  on the HS-Rr  and HS-R2 samples were nonadherent.  Cracks were  also observed within the depleted layer of  the HS-F sample  and at  the  interface  between HfO2  and the SiC-depleted layer  (Fig. 3a).  In the HS-R sam ple,  fracture is observed between the depleted layer and  the HfO2  layer  at  the  center of  the  sample  (Fig. 3b),  whereas  adherent oxide  scales  exist near  the  end of  the  reduced area.  Temperatures Beyond 2000°C  The maximum temperature of  the direct  resistance  test  is  limited only by the available power and the sur vivability of  the  sample. A sample  (HS-Rf) was heated  to failure, where failure was defined as  the sample frac turing such that the electrical path was disrupted. The maximum observed temperature was 2325°C. A micro graph of  the cross  section of  the HS-Rf  sample (Fig. 6)  reveals extensive internal damage. Large pores are found  inside the sample whereas an oxide scale covers  the sur face. The bulk unoxidized material  from the  center of  the  sample  (inset Fig. 6) was  confirmed by EDS to be  SiC and HfB2. The microstructure  suggests  formation  of  a  liquid phase, which is 2347°C in  consistent with the calcu-SiC system.16  lated  eutectic  at  the HfB2  The oxide scale (inset Fig. 6) is composed of HfO2 penetrated by SiO2. Meng et al.13 similarly showed the -SiC sample failure of a ZrB2 at temperatures above 2300°C (2207°C eutectic temperature16), but did not  show any micrographs of  the interior microstructure.  Discussion  The  direct  comparison  of  the  zirconia-resistance -SiC  heated and direct electrical resistance-heated HfB2 2000°C provide  samples  at  insight  to  the  limitations  of  furnace  heating. Due  to  slower  heating  rates  and  contamination  from contact  between  the  sample  and  (a)  (b)  (c)  (d)  Fig. 4.  (a) Micrograph showing the different phases  found in the HS-F oxide scale (1) SiO2, (2) Si-Al-O, (3) HfSiO4 with Ca, and found in (a); (c) Micrograph showing the HfO2 (light) and Si-O-Al (dark) phases found the Si-Al-O phase corresponding to (c).  (4) HfO2; in the HS-R oxide scale;  (b) EDS corresponding to the phases  (d) representative EDS of  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  297  \\x0c', '298  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  (a)  (b)  (c)  Fig. 5.  (a) Heating profile of  sample HS-Rr.  (b) Micrograph of  Fig. 6.  Micrograph of the HS-Rf sample heated to 2325°C. The white-outlined inset shows HfB2 grains (labeled) and the eutectic SiC (dark)-HfB2 (light) structure found in the interior of the sample. The black-outlined inset is the oxide scale composed of HfO2 and SiO2 found on the exterior of the HS-Rf sample.  the oxide found in hottest region of the oxidized HS-Rr sample heated to 2000°C for 1 min two times. oxide found in hottest region of the oxidized HS-R2 sample heated to 2000°C for two minutes. The layers are the same as those found in the HS-R sample: (I) HfO2 penetrated by SiO2, (II) porous HfO2, and (III) depleted HfB2.  (c) Micrograph of  the  crucible,  the HS-F total oxide  scale  thickness  is greater  Fig. 7.  Micrograph showing a Si-Al-O impurity phase in the the HfB2-15 vol% SiC sample. The C signal in the EDS (inset) is from the carbon coating applied to the sample.  bulk of  than that observed in the HS-R sample. The difference  incongruently melting  silicate  and  thus  require  solid in  heating  rates  can  also  explain  the  observation  of  state  diffusion  to  form the  an  extra  kinetic  limitation on its  silicate phase adding formation.17 The  rapid  heating and cooling rates of  the HS-F sample presum ably do not allow for  the separation of glasses with dif HfSiO4  in  the HS-F  sample  but  not  sample. HfSiO4  is  only  stable  below  in the HS-R ~1726°C,17,18  therefore;  its  formation  in  the HS-F  sample  could  occur  during  slow cooling. HfO2  and  SiO2  form an  ferent viscosities or  for  the formation of HfSiO4.  \\x0c', 'In addition,  the  lack of Ca  impurity  in the  resis tance-heated sample suggests  the source of  the impurity  to  be  the CaO-stabilized  zirconia  crucible  or  zirconia  heating  element, whereas  the  presence  of Al  in  both  samples  implies  that  it  is  an inherent  impurity in the -SiC sample  starting powders. For  comparison,  a HfB2  was  heated  in  the  zirconia-resistance  furnace  using  a  Y2O3-stabilized  zirconia  crucible. The  glass  near  the  contact  region of  the sample and crucible was  found to  contain Si, Al, Ca, Y,  and O. Since  the  crucible was  reported to only contain 0.001% Ca,  the zirconia sam ple stand (Part A in Fig. 1) was  the likely source of Ca  in this  test. The HfB2  and SiC powders  are  reported  by the manufacturer  to contain 0.03% and 0.01% Al,  respectively. Figure 7 is a micrograph showing the SiC  grains with a pocket of  impurities  in the  as-processed  material. These areas can be found throughout  the sam ple  adjacent  to  SiC grains  and  are  shown  by  EDS  (inset)  to contain Si, Al, and O. The slow heating rates  and  contact  contamination  issues  of  the  zirconia  ele ment  furnace are not expected in hypersonic flight con ditions  and serve  to complicate  the  analysis of UHTC  oxidation resistance testing. -SiC sample is heated by direct elec When a HfB2  trical  resistance  through multiple heating  and  cooling  cycles,  spallation of  the oxide  scale is -HfO2  suggested by the  presence of  the  repeating SiO2  layers. Such lay ered oxide  structures have not  been reported for  fur nace-heated samples and was not observed in a sample  heated for  the same time (HS-R2) with a single heating  and cooling cycle. There  are  two sources of  stress dur ing oxidation that may lead to fracture during tempera ture  changes:  (i)  thermal  expansion mismatch and (ii)  volume  changes  associated with phase  transformations.  The  coefficient of  thermal  expansion (CTE)  of HfO2  depends on the impurity content and phase, but \\x006 K values are 5 9 10 to 7 9 10 for room temperature to 1250°C with purer HfO2 having lower values.19,20 Gasch et al.21 measured  typical  \\x006  \\x001  the CTE  of  pure  HfB2, pure SiC,  and a  combination of HfB2-20 vol%  SiC to find that \\x006 K the CTE of HfB2-20 vol% SiC was ~5 9 10 ~7 9 \\x006 K \\x001 at 1600°C and fell between the CTE values at room temperature and  \\x001  10  of pure HfB2  (higher)  and SiC (lower)  as  expected by  the rule of mixtures. The transformation of HfO2 from  monoclinic to tetragonal sion at 1642°C) or tetragonal to monoclinic cooling (10% conversion at 1710°C),18,22,23 3-3.5% volume  upon  heating  (10% conver during  is  accom panied  by  a  contraction/expansion  upon  heating/cooling.22,24  This  volume  expansion  could lead spallation of  the oxide scale.  As  the  absolute CTE and modulus of  the multi phase oxide  scale  are not known at  elevated tempera tures,  the main contributing  factor  to oxide  spallation  cannot  be  identified  definitively. However,  if  it  is  assumed that  the  volume  expansion upon phase  trans formation  is  isotropic  then  at minimum the  linear  expansion due  to phase  transformation would be 1%.  To achieve 2000°C to  greater  than  1% linear  expansion  from  room temperature when compared  to  the  bulk, the difference in CTE of the oxide scale and bulk \\x006 K \\x001. The would need to surpass ~4 9 10 reported -SiC and HfO2  range of CTE values for the bulk HfB2 ~2 9 10 \\x006 K  allow  for  a  \\x001  difference  between  the  CTE values, but  the difference could increase at higher  temperatures. Therefore,  it  is  possible  that  the  phase  transformation and CTE mismatch both contribute  to  spallation  of  the  oxide  scale. The  role  of CTE mis match and HfO2 phase  transformation on oxide  scale  adherence deserve further  study.  The limitation of the resistance heater was explored the sample was heated to failure above 2300°C. The  as  entire  sample was  soaked  at  the  elevated  temperature -SiC liquid  allowing for the formation phase inside the HS-Rf  of  the HfB2  sample. Furthermore,  the tem perature may  be  greater  in  the  interior  because  the  oxide scale will not be electrically conductive and is an  effective  thermal  insulator. Under  flight  conditions,  only  the outer  regions of  the  sample would be heated  and  the  high  thermal  conductivity  of  the  diboride  phase would  lead  to  a  temperature  gradient  through  the thickness of  the component. A temperature gradient  is  experienced along  the  length of  the direct  electrical  resistance  sample and can provide  insight  to oxide and  bulk microstructures over a temperature range.  Conclusion  Temperatures up to 2000°C can be  achieved in a  laboratory furnace; however,  these tests  suffer  from slow  heating profiles  and potential  interactions between fur nace materials and the sample being tested. The obser vation of Ca and HfSiO4  in the oxide  scale affects  the  glass  properties,  but  this  is  not  expected  in  a  flight  environment. The use of  resistance heating allows non contact  testing with a high heating profile. Features  like  fracture between the oxide  scale and the depleted layer  www.ceramics.org/ACT  Separating Test Artifacts  in UHTC Oxidation  299  \\x0c', 'and Al  impurities  are  universal  observations  between  both heating tests  and require  further  investigation.  In  addition,  research to stabilize the tetragonal  transforma tion may aid in a more adherent  scale. Resistance heat ing may be  further utilized to  study multiple heating  profiles  and  test materials  for  scale  adherence.  The  resistance  testing  is  limited by  the uniform heating of  the  sample  that would not be  expected in a  real flight  environment. Further comparison of  test methods  such  as  laser heating, oxyacetylene  torch testing, or  scramjet  testing would  be  beneficial  to  understanding material  properties.  Acknowledgments  The authors  thank David Hart, AFRL Air Vehicles  Directorate,  for his assistance with the resistance heater  operation  and  Sindhura  Gangireddy,  University  of  Michigan,  for  the discussion involving microstructures -SiC to  of  resistively  heated  ZrB2  parallel  our  own  observations.  References  1.  J. R. Fenter, “Refractory Diborides 2 1-15 (1971).  as Engineering Materials,” SAMPE Q,  2.  J. W. Hinze, H. C. Tripp, and W. C. Graham, “High-Temperature Oxi dation Behavior of a HfB2 Plus 20 v/o SiC Composite,” Soc., 122 [9] 1249-1254 (1971).  J. Electrochem.  3. W. C. Tripp, H. H. Davis, and H. C. Graham, “Effect of an SiC Addition on the Oxidation of ZrB2,” Ceram. Bull., 52 [8] 612-616 (1973).  4. D. M. Van Wei, D. G. Drewry, D. E. King,  and C. M. Hudson,  “The  Hypersonic Environment: Required Operating Conditions Challenges,” J. Mater. Sci., 39 [19] 5915-5924 (2004).  and Design  5. T. H. Squire and J. Marschall, “Material Property Requirements  for Analy sis and Design of UHTC Components in Hypersonic J. Eur. Ceram. Soc, 30 [11] 2239-2251 (2010).  Applications,”  6. A. Bongiorno,  et al.  “A Perspective  on Modeling Materials  in Extreme  Environments: Oxidation of Ultrahigh-Temperature Ceramics,” MRS Bull., 31 [5] 410-418 (2006).  7. T. Sufang,  J. Deng, S. Wang, W. Liu, and K. Yang,  “Ablation Behaviors  of Ultra-High Temperature Ceramic Composites,” Mater. Sci. Eng. A, 465 [1-2] 1-7 (2007).  8. E. L. Corral  and L.  S. Walker,  “Improved Ablation Resistance  of C-C  Composites Using Zirconium Diboride Ceram. Soc, 30 [11] 2357-2364 (2010).  and  Boron Carbide,”  J.  Eur.  9.  J. Han, P. Hu, X. Zhang, S. Meng, ZrB2-SiC Composites  and W. Han,  “Oxidation-Resistant 799-806  at  2200  °C,” Comp.  Sci. Technol.,  68  (2008).  10. D. D.  Jayaseelan, H.  Jackson, E. Eakins, P. Brown, and W. E. Lee, “Laser  Modified Microstructures in ZrB2, ZrB2/SiC and ZrC,” Soc., 30 [11) 2279-2288 (2010).  J. Eur. Ceram.  11.  S. N. Karlsdottir and J. W. Halloran, “Rapid Oxidation Characterization of Ultra-High Temperature Ceramics,” J. Am. Ceram. Soc., 90 [10] 3233-  3238 (2007).  12. Z. Wang, Z. Wu, and G. Shi, “The Oxidation Behaviors of ZrB2-SiCZrC Ceramic,” Solid State Sci., 13 [3] 534-538 (2010).  13.  S. Meng, C. Liu, G. Liu, G. Bai, C. Xu,  and W. Xie,  “Mechanisms of  Material Failure for Fast Heating up at the Center of Ultra High Temperature Ceramic,” Solid State Sci., 12 [4] 527-531 (2010).  14. T. A. Parthasarathy, M. D. Petry, G.  Jefferson, M. K. Cinibulk, T. Ma thur, and M. R. Gruber, “Development of a Test  to Evaluate Aerothermal  Response of Materials  to Hypersonic Flow Using a Scramjet Wind Tun nel,” Int. J. Appl. Ceram. Technol., 8 [4] 832-847 (2011).  15. C. M. Carney, “Oxidation Resistance of Hafnium Diboride-Silicon Carbide From 1400 to 2000°C,” J. Mater. Sci., 44 [20] 5673-5681 (2009).  16.  L. Kaufman,  “Calculation of multicomponent  refractory  composite phase  diagrams,” NSWC TR 86-242 1 June 1986.  17.  J.-H. Lee, “Ternary Phase Analysis of Interfacial Silicates Grown in HfOx/ Si and Hf/SiO2/Si Systems,” Thin Solid Films, 472 [1-2] 317-322 (2005).  18.  S. V. Ushakov, A. Navrotsky, Y. Yang, S. Stemmer, K. Kukli, M. Ritali  et al.,  “Crystallization in Hafnia and Zirconia-Based Systems,” Phys. Sta tus. Solidi, 241 [10] 2268 (2004).  19.  S. L. Dole, O. Hunter, F. W. Calderwood,  and D.  J. Bray, “Microcrack[11-12] 486-490  ing  of Monolithic HfO2,”  J.  Am.  Ceram.  Soc,  61  (1978).  20.  S. R. Skaggs, “Zero and Low Coefficient of Thermal Expansion Polycrys talline Oxides,” LA-6918-MS September 1977.  21. M. Gasch, D. Ellerby, E.  Irby, S. Beckman, M. Gusman, and S.  Johnson,  “Processing, Properties  and Arc  Jet Oxidation of Hafnium Diboride/Sili con Carbide Ultra High Temperature Ceramics,” 5925-5937 (2004).  J. Mater. Sci., 39 [19]  22. X. Luo, W. Zhou,  S. V. Ushakov, A. Navrotsky,  and A. A. Demkov,  “Monoclinic  to Tetragonal Transformations  in Hafnia  and Zirconia: A  Combined Calorimetric  and Density Functional Study,” Phys. Rev. B, 80  134119 (2009).  23. G. M. Wolten, “Diffusionless Phase Transformations nia,” J. Am. Ceram. Soc., 46 [9] 418-422 (1963).  in Zirconia and Haf 24.  L. Kaufman, E. V. Clougherty,  and J. B. Berkowitz-Mattuck,  “Oxidation  Characteristics of Hafnium and Zirconium Diboride,” Trans. Metall. Soc. AIME, 239 458-466 (1967).  300  International Journal of Applied Ceramic Technology—Carney, Parthasarathy, and Cinibulk  Vol. 10, No. 2, 2013  \\x0c']"
},{
  "_id": 239,
  "PDF": "SiC Depletion in ZrB2–30 vol_ SiC at Ultrahigh Temperatures.pdf",
  "Text": "['SiC Depletion in ZrB2  -30 vol% SiC at Ultrahigh Temperatures  Kathleen Shugart,  ‡,§,*,†  and Elizabeth Opila  ‡,*  ‡  Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22904  §  UES Inc., Dayton, Ohio  The  formation of a porous SiC-depleted  region in ZrB2-SiC due to active oxidation at ultrahigh temperatures was charac terized. The presence/absence of SiC depletion was determined (1300°C-1800°C) temperatures T < 1627°C,  at a series of (5 min-100 h).  and  times  At  SiC depletion was not a ZrO2 + C/borosilicate oxidation product layer sequence was observed above the ZrB2-SiC base material. At T ≥ 1627°C, SiC was depleted in the ZrB2 matrix below the ZrO2 and borosilicate oxidation products. The SiC depletion was attributed to active oxidation  observed.  Instead,  the  formation  of  of SiC to form SiO(g). The transition between C formation in (T < 1627°C) and SiC depletion in ZrB2 (T ≥ 1627°C) ZrO2 is attributed to variation in the temperature  dependence  of  thermodynamically  favored  product  assemblage  influenced  by  the local microstructural phase distribution. The growth kinet ics of  the SiC depletion region is  consistent with a gas-phase  diffusion-controlled process.  I.  Introduction  Z RB2-SIC is proposed for sharp leading edges of Sharper wing leading edges and noses allow speeds.1,2 When  thermal protection systems  for  advanced  hypersonic  vehicles.  for  better  maneuverability  and  higher  the  radius  of  the wing  is made  smaller  and  as  speed  is  increased,  the  surface  temperatures  seen by the  spacecraft upon reentry to 1700°C-3000°C due to surface.2-4 operating  the Earth’s  atmosphere  increase  to  closer  proximity  of  shockwaves  to  the  vehicle  Such  high  temperatures  are  beyond  the  capabilities  of  current  thermal  protection systems (TPS), TPS materials.3-7 are a family of materials in excess of 3000°C, making  necessitating  the  development  of  new  Ultrahigh-temperature  ceramics  that have melting temperatures  them attractive  for TPS aerospace  applications. One mem ber  of  this  family, ZrB2 (and ZrB2-based materials), been extensively studied since the 1960s, due to its favorable melting temperature, 3245°C, 6.09 g/cm3,  has  relatively low theoretical den sity,  strength  retention at high temperatures, to erosion/corrosion.3,4,7-9 for non-oxide TPS materials is  chemical  stability,  and resistance  One of  the  largest  concerns  oxidation  behavior  at  ultrahigh  temperatures.  SiC is  com monly added to ZrB2 to improve the oxidation resistance at 1100°C.10-12 However, temperatures above resistance of ZrB2-SiC is inadequate for long-term use and poorly understood. The oxidation of ZrB2-SiC occurs via the following reactions to form ZrO2 and borosilicate glass:  the  oxidation  ZrB2 ðsÞ þ 5 2  O2 ðgÞ ¼ ZrO2 ðsÞ þ B2O3 ðl or gÞ  (1)  SiCðsÞ þ 3 2  O2 ðgÞ ¼ SiO2 ðsÞ þ COðgÞ  (2a)  SiC(s) þ O2 ðgÞ ¼ SiOðgÞ þ COðgÞ  (2b)  xSiO2 ðs,lÞ þ yB2O3 ðlÞ ¼ xSiO2 \\x01 yB2O3 ðs,lÞ  (3)  The  oxidation  of ZrB2-SiC at 1500°C is tions (1) and (2a)] and forms a protective layer of borosilicate glass above ZrO2, along with CO(g).13,14 The protective capability of the borosilicate layer is likely to vary consider passive  [Reac ably with temperature due to compositional layer as boria volatilizes.15 The  changes  in the  borosilicate  transition from  passive oxidation [Reaction (2a)]  to active oxidation [Reac tion (2b)] of SiC occurs at high temperatures and low oxygen partial pressures.4,6,10,12,16-22 The oxygen partial pressure at the ZrB2/ZrO2 interface is sufficiently low such that at high temperatures the SiC is expected to undergo active oxidation depletion.10,23  [Reaction (2b)]  leading  to  SiC  The  gases  formed at  low oxygen partial pressures at  the substrate/oxide  interface  in  turn  oxidize  at  the  higher  oxygen  partial  pressures  found near  the oxide/gas  interface  [Reactions (4)  and (5)].  SiOðgÞ þ 1 2  O2 ðgÞ ¼ SiO2 ðgÞ  (4)  COðgÞ þ 1 2  O2 ðgÞ ¼ CO2 ðgÞ  (5)  Conditions  for  formation of  the SiC depletion layer,  for  example,  the  transition from passive  to active oxidation in  the presence of ZrB2, containing 10 vol% or  are not well  characterized. Materials  less of SiC have not  shown a deple tion layer, which may be due to the lack of interconnectivity the SiC grains.16,22,24 Connectivity of SiC grains may be  of  necessary to provide a continuous path for  inward transport  of oxidizing species and outward flow of  reaction products.  Under  the  low oxygen potential at  the ZrB2/ZrO2 to form SiO(g) and CO(g) which dif interface,  the interior SiC reacts  fuse  outwards. Continuous  pores  are  left  behind  through  which O2(g) diffuses inwards, the material. The SiC depletion layer has been reported in ZrB2-SiC based materials containing more than 10 vol% SiC when they are oxidized at temperatures as low as 1400°C,10,19 greater than 1600°C,16 1900°C,12 or not at all.7,18  further depleting the interior of  only  lower  than  In addition,  the thermodynamics  and kinetics for the depletion reaction are poorly understood.  A thermodynamic model based on superposition of volatility  diagrams  for ZrB2 and SiC has been proposed by Fahren B. Fahrenholtz—contributing editor  Manuscript No. 35400. Received August 1, 2014; approved January 21, 2015.  *Member, The American Ceramic Society.  †  Author to whom correspondence should be addressed. e-mail: kns9a@virginia.edu  1673  J. Am. Ceram. Soc., 98 [5] 1673-1683 (2015)  DOI: 10.1111/jace.13519  © 2015 The American Ceramic Society  Journal  \\x0c', 'holtz  to elucidate  the SiC depletion behavior,10  though its  prediction of onset  temperature is not  in agreement with the  results in this work.  The porous  structure  formed during SiC depletion could  lead to reduced load-bearing capability or reduced thermal conductivity of ZrB2-SiC for applications at ultrahigh temperatures. Therefore, understanding conditions for depletion  layer formation are necessary future use.16 The objectives of  to  develop  this material  for  this work were to (1)  identify  temperature and time conditions for formation of SiC depletion in ZrB2-30 vol% SiC and (2) develop a thermodynamic and kinetics-based understanding of this depletion mecha nism to enable life prediction.  II.  Experimental Procedure  ZrB2-30 vol% SiC bars were provided by Dr. Fahrenholtz (Missouri University of Science and Technology).19,25,26 The  specimens  were fabricated using attrition-milled powders then hot pressed. WC contamination (~2 wt%)  which were  was observed due  to the  attrition milling, which used WC  milling media in polyethylene jars. 40 mm 9 4 mm 9 3 mm dimensions were  Bars  of  cut  from a series  of billets using an automated surface grinder and were fin ished using 1200 grit diamond abrasive. Two specimen con figurations were used for oxidation testing. Specimens were from bars to size (~7 mm 9 4 mm 9 3 mm) using a dia cut  mond blade, as  shown in Fig. 1(a),  for box furnace  testing.  A second configuration of bridge specimens were machined 15 mm 9 3 mm 9 1.5 mm  to  dimensions  of  by  Bomas  Machining Specialties  (Somerville, MA). The middle 3.5 mm  was  thinned to 0.5 mm for  resistive heating oxidation tests  [Fig. 1(b)]. Before  heating,  all  specimens were  cleaned  in  detergent and DI water, DI water, acetone, and then etha nol. Tests between 1300°C and 1550°C were box furnace with molybdenum disilicide  carried out  in a  heating  elements  (RapidTemp; CM Furnaces Inc., Bloomfield, NJ) under stag nant ambient air conditions. This furnace permitted the spec imen to be  inserted and removed while  the  furnace was at  temperature, allowing for precise control of  the thermal his tory for  each exposure. Specimen weights and surface area  measurements were acquired before and after  each box fur nace  exposure. The  specimens were held in an yttria-stabi lized zirconia (YSZ; Ortech Inc., Sacramento, CA) boat  for  furnace  testing, as  shown in Fig. 1(a). This boat prevented  contamination of  the specimen and allowed for easy insertion  to and removal from the furnace. The oxidation behavior of ZrB2-30 vol% SiC was tested at ultrahigh temperatures (1500°C-2000°C), using a resistive heating method first developed by Karlsdottir and Halloran, which they have called “the ribbon method.”17 This method  resistively heats ~100 Amps, specimen. As \\x005Ω\\x01cm, the tance of ZrB2 is similar to metals, ~1 9 10 approach works well.27 The machined “ribbon” or bridge  the  specimen  using  a  high  current,  that  runs  through the  resis this  specimen, as  seen in Fig. 1(b), was used for  this  technique.  The  thin  bridge  region  allows  the  center  of  the  specimen,  which was not  in direct contact with any surfaces,  to heat  to  ultrahigh temperatures. The  supporting ends  remained rela tively cool, minimizing high-temperature  contamination and  interaction of  the specimen with the heating apparatus. The  resistive  heating method  also  has  the  advantages  of  being  more  economical  than  an  arc-jet  facility  and  allowing  for  more control of the temperature than using an oxyacetylene torch, both alternative techniques for reaching T > 1550°C.  The specimen temperature was measured using an emissiv ity-correcting infrared pyrometer with a spot  size of 1.5 mm  (Pyrofiber Lab, Pyrometer  Instrument Company, Windsor,  NJ), which automatically corrected for the changing emissiv ity  of  the  oxidizing  surface  using  pulsed  laser  technology.  The pyrometer also controlled the current which ran through  the specimen by providing the temperature reading to a con trol unit with a Eurotherm controller (BPAN controller; Mi cropyretics Heaters  International  Inc., Cincinnati, OH). The  control unit provided the current  to maintain the desired set  temperature. The  current passed through a transformer to (~90 Amps). The  step the  current up to the  required level  current  running through the  specimen was monitored using  an AC current  transducer and continuously recorded using a  NI-6009 data acquisition card (DAQ) and both the tempera ture and the current were recorded by a computer. The resis tive heating system is shown schematically in Fig. 2.  For  testing,  each end of  the bridge  specimen was placed  on flattened copper wire with 1.05 mm 9 2.85 mm, which was  a  cross  section  of  attached  to  the  current  feedthroughs  in the  specimen chamber  (Fig. 3). A piece of  platinum foil (0.06 mm thick; Heraeus Materials Technology,  Chandler, AZ) was placed between the copper wire and the  specimen  to minimize  reactions  between  the  specimen  and  the wire. The  specimen was  held  in  place  using  toothless  alligator  clips,  which  were  electrically  isolated  using  0.8-mm-thick  sheets  of Maycor  (a  glass-mica  ceramic,  McMaster-Carr, Robbinsville, NJ) or  stabilized ZrO2 Inc, Houston, TX). A bar of stabilized  (Cera mic Technologies,  ZrO2 was super glued to the side of stability against torquing when clamping. B2O3(g) condensation on the window was prevented by flowing O2(g) through the chamber at a rate of 900sccm. The O2(g) flowed through drierite (CaSO4) prior to entering the resistive heating chamber to maintain low humidity. Pt (/0.508 mm; Sigmund Cohn, Mount Vernon, NY) and (/0.203 mm; Alfa Aesar, Ward Hill, MA) wires were (Tm = 1555°C and 1768°C) readings. The wires of each metal  the specimen to provide  Pd  used as melting point  standards  to confirm the pyrometer  (a)  (b)  Fig. 1.  ZrB2-30 vol% SiC specimen (a) (YSZ) boat for box furnace oxidation exposure  in yttria-stabilized zirconia  (scale  is  in cm)  (b)  machined into bridge shape for resistive heating.  1674  Journal of the American Ceramic Society—Shugart and Opila  Vol. 98, No. 5  \\x0c', 'were wrapped tightly around preoxidized bridge specimens to  maintain good contact with the thin bridge region. The speci mens were  then  inserted  into  the  resistive  heating  system.  The temperature of the specimen was slowly increased (10°C-20°C/min) while the wire was observed. When the wire melted, the temperature reading of the pyrometer was  recorded. Observed melting temperatures of the Pd and Pt wires were 1527°C and 1752°C, respectively. These wire melting tests indicate the pyrometer reads within 28° of the tabu lated melting temperature interest.28,29 Table I  in  the  temperature  range  of  lists  the conditions and heating method  used for each oxidation test.  Specimens were prepared for Scanning Electron Micros copy  (SEM;  JEOL 6700F, Tokyo,  Japan;  or FEI Quanta  600F, Hillsboro, OR)  and Energy Dispersive  Spectroscopy  (EDS; Princeton Gamma-Tech Inc., Princeton, NJ; Oxford Instruments Aztec X-MaxN 150, Concord, MA)  or  as  follows. Box-furnace-oxidized  specimens were mounted  in  epoxy while under  low vacuum. After curing,  the specimens  were cross-sectioned using a diamond saw, and remounted in  epoxy. The first  epoxy step was performed to minimize  the  oxide spallation during sectioning and polishing. The cross sections were polished using diamond grit down to 1 l and a  mixture of  ethylene glycol and 200 proof  ethanol. Ethylene  glycol and 200 proof ethanol were used as lubrication during  polishing to minimize loss of B2O3, which reacts readily with water. Even with careful polishing SiO2/C/SiC pull out from the ZrO2 oxide layer was observed, resulting in the appearance of selective depletion, as will be discussed below. After  resistive  heating  oxidation,  the  bridge  specimens  were  snapped in half  in the middle of  the bridge  region, which  often  occurred  naturally  upon  cooling. The  fracture  cross  sections were then characterized.  Specimens were coated with carbon prior  to SEM charac terization to provide  a  conductive  layer  (Precision Etching  and Coating  System). The  thicknesses  of  the  borosilicate,  ZrO2 and (if present) SiC depletion layers were determined by measuring the layer thickness in a minimum of 20 loca tions over  the  cross  section and averaging. SEM character ization was performed using 5 kV to maximize EDS intensity  for light elements.  Thermodynamic  calculations were performed using Fact Sage, a computer program designed for chemical calculations.30 The Fact  thermody namics  Pure  Substances  database  was utilized.  III.  Experimental Results  (1)  Comparison of Box Furnace to Resistive Heating  Concerns  exist  that  internal heating by resistive heating can  result  in different oxidation rates and microstructures when  compared with conventional box  furnace  radiative heating.  The validity of  the resistive heating tests was established by  comparing specimens  exposed in the box furnace and resis tive  heating  system. Both  tests were  performed  on  bridge  specimens of 1500°C (near  ZrB2-30 vol% SiC the maximum temperature  oxidized  for  20 min  at  capability  of  the  box  furnace.) The  specimens were  exposed in stagnant  air  Fig. 2.  Schematic of resistive heating system.  Fig. 3.  ZrB2-30 vol% SiC bridge in resistive heating system.  specimen prepared for oxidation  May 2015  SiC Depletion in ZrB2-SiC  1675  \\x0c', '(box furnace) and flowing O2 mens were fractured and characterized using SEM. The oxide  (resistive heating). Both speci cross  sections appeared similar, as presented in Fig. 4. The  box furnace 5.5\\x060.8 lm and specimen showed an average borosilicate 13.1\\x061.6 lm thickZrO2 the mean and standard deviation were determined  ness  of  thickness  of  where  showed an average borosilicate thickness of 7.1\\x060.8 lm and from 23 measurements. The resistive heating specimen ZrO2 thickness of 11.0\\x062.1 lm where the mean and standard deviation were determined from 28 measurements. A t-test  was performed comparing the mean oxide  layer  thicknesses  formed in the different  testing environments. The mean oxide  thicknesses  (both glass  layer and ZrO2 statistically significantly different at  layer)  formed in the  two environments  is  the  95% confidence  level. We do, however,  emphasize  that  the  box furnace studies were conducted in air and the resistively  heated samples were oxidized in oxygen. The higher oxygen  partial pressure  resulted in greater glass formation and the thickness.31 We  corresponding lower ZrO2 ference in oxygen partial  feel  that  the dif pressure  rather  than  the  heating  method is  responsible  for  the differences  in layer  thickness.  We  are most  concerned  that  the  oxide morphology  and  phases present were  consistent with each other: borosilicate of ZrO2 + carbon. The oxide morphology shows that the two tests are equivalent at 1500°C and no significant artifacts of  glass  formed  over  a  layer  similar  resistive heating were  observed.  (2)  Oxide Growth Kinetics  for ZrB2-SiC is Oxidation kinetics often determined using mass gain/surface area versus time.12,21,31 For specimens oxi dized using resistive heating, kinetics  cannot be determined  using mass change, as the entire specimen was not at  temper ature throughout  the test.  In addition,  the borosilicate oxide  layer has inconsistent  thicknesses due to bubbling. Therefore,  oxidation growth kinetics was determined using ZrO2 oxide thickness measurements. Results are given in Table II. A found in reference.31  comparison to literature values  can be  While large variation is seen for the ZrO2 growth rates at all temperatures31 leading to uncertainty in the reported parabolic growth rate constants, kp (as represented by the low R2 values for fit to the data), an order of magnitude increase in temperatures of 1650°C  the kp calculated for oxides grown at  Table I.  Heating Method for Each Oxidation Test  Time  (min)  Temperature (°C)  Heating  method  No of  specimens  tested under  identical  conditions  No SiC  depletion  10  1300  Furnace  1  10  1300  Resistive  1  30  1300  Furnace  2  50  1300  Furnace  2  10  1500  Furnace  2  10  1500  Resistive  3  20  1500  Resistive  1  30  1500  Furnace  4  50  1500  Furnace  2  100  1500  Furnace  11  6000  1500  Furnace  1  10  1550  Furnace  3  30  1550  Furnace  3  50  1550  Furnace  3  55  1600  Resistive  1  40  1627  Resistive  1  60  1627  Resistive  1  Start of  SiC depletion  70  1627  Resistive  1  10  1650  Resistive  2  20  1650  Resistive  1  Distinct  layer  of SiC depletion  35  1650  Resistive  1  45  1650  Resistive  1  50  1650  Resistive  1  55  1650  Resistive  1  60  1650  Resistive  1  70  1650  Resistive  1  10  1700  Resistive  1  15  1700  Resistive  1  20  1700  Resistive  2  25  1700  Resistive  1  30  1700  Resistive  2  35  1700  Resistive  2  40  1700  Resistive  2  5  1800  Resistive  1  10  1800  Resistive  2  15  1800  Resistive  2  20  1800  Resistive  1  (a)  (b)  Fig. 4.  ZrB2-30 vol% SiC oxidized for 20 min at 1500°C using (a) standard box furnace in stagnant air (b) resistive heating system in flowing O2.  Table II.  Experimental Results for Oxidation Kinetics of ZrB2-30 vol% SiC  Temperature (°C)  Mass gain kp (mg2/cm4 h)  ZrO2 growth kp (lm2/h)  R2 for ZrO2  growth kp fit  line  1300  4.0  295  0.91  1400  7.4  280  0.61  1500  4.4  390  0.94  1650  8500  0.21  1700  28 000  0.21  1800  42 000  0.52  1676  Journal of the American Ceramic Society—Shugart and Opila  Vol. 98, No. 5  \\x0c', 'and above is observed compared to oxides grown at temperatures of 1500°C and below. This change in kp corresponds to observed changes in the microstructure, as discussed below.  (3)  Oxide Morphology Regime I: Temperatures <1627°C  At  temperatures below 1627°C,  specimens  showed two dis tinct  layers  after  oxidation. The  top  layer  of  borosilicate  glass with an underlying layer  consisting of  spherical ZrO2 grains was observed as shown in Fig. 5 and 6. At early times in the oxidation, such as 1500°C for 20 min, both SiC and C  grains  were  interspersed  between  the  ZrO2 temperature was  grains  as  in  Fig. 5. However, as  time and/or  increased,  minimal Si remained between the ZrO2 grains, though C is clearly present, as in Fig. 6 (exposure at 1600°C for 55 min).  It  is  important  to note  that no porous  layer  forms  in the  ZrB2 base material due to oxidation at temperatures below 1627°C for ZrB2-30 vol% SiC, whether oxidized in the box furnace or resistive heating system.  (4)  Oxide Morphology Regime II: Temperatures ≥1627°C  At  temperatures  of  1627°C and  above,  specimens  showed  three distinct  layers after oxidation. A top layer of borosili cate  glass was  observed  over  a middle  layer  consisting  of  columnar  ZrO2 Beneath these two layers was a third layer which consisted of  grains  filled  in  with  borosilicate  glass.  ZrB2 grains and porosity resulting from SiC depletion. Figure 7 is an EDS map of a specimen oxidized at 1650°C for  50 min, which clearly shows a region containing only Zr and  B, with no Si or C signal, beneath the O containing layers.  The necessary time-temperature conditions for SiC deplein ZrB2-30 vol% SiC have been mapped in tion to occur temperatures below ~1627°C, no SiC depletion  Fig. 8. For  formed  regardless  of  the  time  of  oxidation  or  oxidation  method (box furnace or  resistive heating system.) Maximum  exposure 1600°C,  times  were  100 h  and  1500°C 55 min at and increased (1650°C,  respectively. As  the  temperature  45 min),  partial  depletion was  observed where  some  SiC  grains were partially  consumed (Fig. 9)  leaving void space. 1650°C that C  There was  indication  at  10  and 20 min  at  grains depleted in Si and SiC depletion coexisted, as  seen in  Fig. 10. These are the only two test conditions which resulted  in both types of SiC oxidation and captured the microstruc tural dependence of the transition temperature.  (5)  SiC Depletion Growth Kinetics  Figure 11 plots the depletion layer thickness versus oxidation  time for all  tests. The incubation time for depletion of SiC to  form porosity within a layer decreased with increasing tem perature. The depletion layer  clearly grew over  time at  the  base material/depletion layer interface. Due to the significant  scatter  in the time dependence for depletion layer growth, a  clear rate law is not obvious. The depletion layer growth was  therefore fit  to both linear and parabolic rate laws  to exam ine  the  temperature dependence. These  results are  shown in  Fig. 12 and discussed further below. The top of  the depletion  layer  shows  some SiO2 below the depletion layer/ZrO2 indicating that oxidation of the SiO(g) to SiO2 (Fig. 13), as previously  inter face,  likely  occurs near the ZrB2/ZrO2 interface reported4,10,12 and discussed below.  IV.  Discussion  (1)  Temperature Considerations  The use of an emissivity-correcting pyrometer and the rapid  oxidation kinetics should lead to steady-state emissivities and the temperature.28,29 However, fluc accurate measurement of  tuations  in  the  temperature  reading  of  the  pyrometer  occurred during 53°C. This  testing with  standard  deviations  of  5°C-  range of  standard deviations  is  consistent with  Fig. 5. EDS indicating the presence of both C and SiC particles 1500°C for 20 min in flowing O2 using resistive heating.  in the ZrO2  layer of a ZrB2-30 vol% SiC fracture section after oxidation at  May 2015  SiC Depletion in ZrB2-SiC  1677  \\x0c', '1678  Journal of the American Ceramic Society—Shugart and Opila  Vol. 98, No. 5  Fig. 6.  Secondary electron image and EDS maps of ZrB2-30 vol% SiC fracture using resistive heating. C between ZrO2 grains is clear.  section after oxidation at 1600°C for 55 min in flowing O2  Fig. 7.  Secondary electron image and EDS maps of ZrB2-30 vol% SiC fracture using resistive heating. SiC depletion under ZrO2 and borosilicate layers is clear.  section after oxidation at 1650°C for 50 min in flowing O2  calibration tests which differed by \\x0028°C the melting point and \\x0016°C from the melting points of Pt and Pd, respectively. The fluctuating temperature readings are attributed to  bubbling of  the borosilicate which makes  temperature mea surement difficult. The fluctuation may also be due in part  to  difficulties  in maintaining  the  needed  current  for  constant  of  temperature due to the continuously changing cross conductive material. Reporting of 1627°C as temperature for change in the oxidation mechanism is uncertain by as much as 55°C. While a transition temperature  section  critical  the  is  not reported in the literature, these results are consistent with microstructures published by other researchers.19,21,32-34  (2)  Oxide Morphology Regime I: Temperatures <1627°C  the ZrO2 + C oxide layer relative Due to the low hardness of to the ZrB2 + SiC layer, pullout of the C between the ZrO2 grains during polishing can be easily mistaken for porosity.  In addition,  the presence of C is difficult  to distinguish from  epoxy for specimens mounted and impregnated for cross-sec tional characterization. Fracture sections avoid both compli cations,  enabling  the  definitive  identification  of C as  in  Figs. 5, 6, 7, and 10.  It  is possible that  some reports of SiC  depletion  found  in  the  literature  are  actually  due  to  the  decomposition of SiC into carbon and gaseous  species and  not complete removal of SiC.  \\x0c', 'The preferential oxidation of  the Si  from SiC leaving C is  attributed to the  low partial pressure of oxygen under  the  borosilicate scale. Previous work in related systems has dem onstrated  similar  behavior.  Cooper  et al.  observed  the  formation of C due to the oxidation of SiC fibers composites.35,36 They  in glass  ceramic  called  this  phenomena  “car bon-condensed  oxidation”  as  given  by  the  displacement  Reaction (6).  SiC þ O2 ðgÞ ! SiO2 ðgÞ þ CðsÞ  (6)  Their work  suggested  that  at  low oxygen  activities  the  more  stable  oxide,  SiO2, also proposed that continued direct  formed  rather  than CO(g). They  formation of SiO2 adjato the C required diffusion of Si through C. Work by  cent  Katsui et al. has  suggested that oxidation controlled by CO  (g)  diffusion  through SiO2 in both a C layer and bubbling within the SiO2, and a similar mechanism could explain the behavior seen here.37  during  oxidation  of  SiC would  result  However,  according  to Zheng  et al. CO(g) molecules  are  smaller  than O2 and transport of CO(g) oxidation rate.38 In addition, Katsui  is unlikely to limit  the  et al.  showed  no  change in their log(k) vs. 1/T slope, which would be expected  for the transition from O2(g) diffusion limited to CO(g) diffusion limited oxidation as proposed.  We were unable to reproduce the Ellingham Diagram in Fig. 9(B) of Cooper et al.35 using the FactPS database,30 spe cifically  the  transition  temperatures  between Reaction (6)  (low-temperature  regime), Reaction (2a)  (intermediate-tem perature  regime),  and  Reaction (2b)  (high-temperature  regime). We propose the following reaction considering oxidation of ZrB2-SiC at low partial pressures of oxygen.  SiC þ 1 2  O2 ðgÞ ! SiOðgÞ þ CðsÞ  (7)  This new finding of C in the ZrO2 scale indicates CO(g) production below 1627°C is unlikely to form at sufficient pressures to produce bubbles in the borosilicate scale, contrary to our previously proposed model.31 Other gases that could contrib ute to bubble formation are SiO(g) or boiling B2O3. However, SiO(g) will react with inward diffusing oxygen to form con densed phase SiO2. In addition, the boiling temperature for B2O3 at 1 atm pressure is 2058°C (FactSage PS database).30 Thus we have not identified a plausible explanation for the  bubbles  observed  in  the  borosilicate  scale  at  temperatures  below the transition temperature for SiC depletion.  (3)  Transition from C Condensed Oxidation to SiC  Depletion  Reaction (8) represents the observed phases at the base mate<1627°C. Reaction (9) represents the observed phases at the base matetemperatures ≥1627°C. Note  rial/oxide  interface  for  temperatures  Similarly,  rial/depletion layer  interface for  that ZrB2 observed  is only included in Reaction (9)  to represent  the  interface  and  phases  for  comparison  to Reac tion (8), but does not enter into the reaction.  ZrB2 þ SiC þ 3O2 ðgÞ ! ZrO2 þ C þ SiOðgÞ þ B2O3 ðlÞ  (8)  ZrB2 þ SiC þ O2 ðgÞ ! ZrB2 þ SiO þ COðgÞ  (9)  The balanced reactions  for  the  compositions  studied will  be discussed below.  Figure 14  is  an  Ellingham  diagram,  including  Reac tions (8) and (9), plotted using results Sage.30 Note that the slopes of the lines in Fig. 14 reflect the the entropy change in each reaction.39 For Reac calculated with Fact negative of  tion 8 there  is a net  loss of  two gas molecules,  reflecting a  significant decrease in entropy. The free energy of  formation  is  therefore  expected to increase with temperature  (positive  slope). For Reaction (9),  in which there is a net production  of one gas molecule (increase in entropy),  the slope is nega tive. These  slope  differences  necessitate  a  transition  from  Reactions (8)  to (9) at high temperatures as the reaction with  a lower Gibbs energy will be favored.  Reactions (8) and (9) are  rewritten in Table III balanced  for 30 vol% SiC and shown in Fig. 14. Note that  the transi tion temperature  calculated for Reactions (8) 30 vol% SiC occurs at 1725°C, a temperature higher the 1627°C observed experimentally. SiC mole ratio is varied, the net change  and  (9) with  than  If, however,  the ZrB2/ in gas molecules/  mole of O2 varies unaffected. As the mole fraction of SiC increases, the slope of DG for Reaction (8) decreases and the transition temperafrom oxidation products of C + ZrO2 + SiO(g) + B2O3 [Reaction (8)] to SiO(g) + CO(g) [Reaction (9)] moves to  for Reaction (8), whereas Reaction (9)  is  ture  lower temperatures. The computed transition temperature occurs between 1627°C and 1650°C in agreement with the ZrB2-60 vol% thermodynamics of  experimental temperature ZrB2-65 vol% SiC. Thus, rich microstructures better  for  SiC  and  locally  SiC reflect  the  observed  transition  Fig. 9.  Fracture  section of ZrB2-30 vol% SiC oxidized at 1650°C flowing O2 using resistive heating system showing partially removed SiC grains  for  45 min  in  (arrows)  just above  the base material/  SiC depletion layer interface (below view).  in ZrB2-30 vol% Fig. 8. Oxidation conditions for SiC depletion SiC. All graphed exposures T < 1550°C were conducted in stagnant air using the box furnace. All graphed exposures T > 1550°C were  conducted in 900sccm flowing O2 using the resistive heating system.  May 2015  SiC Depletion in ZrB2-SiC  1679  \\x0c', '1680  Journal of the American Ceramic Society—Shugart and Opila  Vol. 98, No. 5  Fig. 10.  electron image and EDS maps of ZrB2-30 vol% SiC fracture 20 min in flowing O2 using resistive heating. Squares mark pores left by SiC depletion. Circles mark C remaining after oxidation of Si.  High magnification secondary  section after oxidation at 1650°C for  Fig. 11.  SiC depletion depth versus  time  in minutes  for oxidized  ZrB2-30 vol% SiC.  Fig. 12.  Log(k) versus 1/T for both linear and parabolic growth of in oxidized ZrB2-30 vol% SiC. kl is lm/min  the SiC depletion layer and kp is l/min1/2.  temperature from SiC depletion to ZrO2 + C oxide layer formation.  To test  this model, we used the FactSage reactions module  to  calculate  the  free  energy  of  reaction  for Reactions (2a)  the  lower  limit  at  at  an  and (2b)  calculated  and calculated the  transition temperature for SiC 1883°C (2156 K) oxidation to SiO2 or SiO(g) to be \\x0013 atm. This oxygen partial pressure of 2.7 9 10 compares favorably with the Si-C-O volatility diagram of Heuer and Lou40 who  SiO2 2100 K to be at an oxygen partial pressure of approximately \\x0012 atm. To 10 further elucidate the transition temperature for ZrB2-SiC from ZrO2 + C formation to SiC depletion we have calculated the T, PO2 conditions at which this occurs for 30, 60, and 65 vol% SiC. The transition temperature and \\x0013 atm; oxygen partial pressures are 1725°C, PO2 = 1.3 9 10 \\x0014 atm; 1633°C, PO2 = 8.5 9 10 and 1618°C, PO2 = 7.8 9 10 \\x0014 atm for ZrB2 with 30, 60, and 65 vol% SiC, respectively. Thus, these higher SiC contents result in  stability  of  more  reducing conditions, with corresponding transitions  to  SiC depletion at  lower temperatures.  The  from ZrO2 + C to SiC depletion corresponds with the order of magnitude  in stable oxidation products  change  jump in ZrO2 oxide thickness, as documented in Table II. As the oxidation mechanism changes from Reactions (8) to (9),  the C oxidizes  to CO(g) generating additional gas products.  The additional gas products  result  in increased bubbling of  the borosilicate glass, which was observed in posttest analysis  (Fig. 15.)  Increased bubbling would reduce  the  thickness of  the protective glass  layer, allowing ingress of oxygen to the  ZrB2, thereby increasing the ZrO2 growth rates as previously noted.22  (4)  The  Oxide Morphology Regime II: Temperatures ≥1627°C  presence  of  partially  oxidized  SiC grains  at  the  base  material/SiC depletion layer  interface  is a strong indication  of  active oxidation (Fig. 9), by Reaction (2b). The PO2  at  \\x0c', 'May 2015  SiC Depletion in ZrB2-SiC  1681  Fig. 13.  SE image and EDS maps of ZrB2-30 vol% SiC fracture heating showing borosilicate glass beneath the ZrO2/ZrB2 interface.  section after oxidation at 1650°C for 50 min in flowing O2 using resistive  Fig. 14.  Ellingham diagram, plotted using results of FactSage calculations,  showing reactions of ZrB2-SiC with O2, balanced for 1 mol of O2  and calculated using increasing quantities of SiC.  Table III.  Balanced Reactions for Transition in ZrB2-SiC Oxidation Behavior  Vol% SiC  8  8  8  9  30  60  65  30  Balanced reaction for 1 mol O2 (mol%)  Transition temperature (°C)  0.35ZrB2(s) + 0.22SiC(s) + O2(g) = 0.35ZrO2(s) + 0.22C(s) + 0.22SiO(g) + 0.35B2O3(l) 0.2865ZrB2(s) + 0.657SiC(s) + O2(g) = 0.2865ZrO2(s) + 0.657C(s) + 0.657SiO(g) + 0.2865B2O3(l) 0.2549ZrB2(s) + 0.725SiC(s) + O2(g) = 0.2549ZrO2(s) + 0.725SiO(g) + 0.2549B2O3(l) 0.725C(s) 1.564ZrB2(s) + SiC(s) + O2(g) = 1.564ZrB2(s) + SiO(g) + CO(g)  +  1725  1630  1610  N/A  Note that ZrB2 is included in reaction (9)  to represent  the observed interface and phases.  \\x0c', 'the ZrO2/ZrB2 thermodynamically  interface must be  low enough that SiO(g)  is  stable  compared to SiO2(s,l). The the ZrO2/ZrB2 \\x0020 atm at interface [Reaction (1)] was 5.3 9 10 1650°C, assuming the activities of all condensed phases are unity. For comparison, Fahrenholtz10 calculated the oxygen partial pressure in between 4.1 9 10 \\x0016 atm 1.8 9 10 interface as  equi librium PO2 calculated to be  at  that  the SiC-depleted layer to \\x0011 Pa, 1.8 9 10 or 4 9 10 1500°C. However, SiO2 shown in Fig. 7 and schematically  lay  \\x0014  and  \\x0019  and  at  is also present near  this  in Fig. 16. Under  such  low partial  pressures, \\x006 atm and O2(g) SiO2 would reduce \\x006 atm by the of 2.7 9 10 reverse of Reaction (4), providing a source of oxygen for further  to  SiO(g)  at  a  pressure 1.1 9 10  at  a  pressure  of  active oxidation of SiC. This  mechanism requires  consideration of  equilibrium thermody namics  at  and adjacent  to the ZrB2/ZrO2/SiO2 the observed microstructures. The O2(g) generated reduction of SiO2 adjacent to the ZrB2/ZrO2 interface through the pores to enable rapid active oxidation  junction to  interpret  by  diffuses  of SiC to greater depths  in the ZrB2 as in Fig. 16. It should be noted that the work of Parthasarathy et al.41 and Holcomb and St. Pierre42 show that a CO-  shown schematically  CO2 oxygen  exchange mechanism must  be  operating  at  these  low  partial  pressures  to  enable  oxygen  transport,  however, we show O2 simplicity.  transport  in the Fig. 16 schematic for  (5)  Depletion Kinetics  Figure 12  shows  that  the  temperature  dependence  of  observed depletion layer growth assuming either a parabolic  or  linear  rate-limiting step.  If  the rate-limiting step involved  bond breaking or solid-state diffusion,  there would be a posi tive exponential  temperature dependence of  the linear or par abolic rate constant, respectively. Gas-phase diffusion, on the other hand has a T3/2 dependence as derived from the kinetic  theory  of  gases. The  assumed  parabolic  rate  constants,  in  particular, do not show a positive temperature dependence in  Fig. 12. The limited data in Fig. 12 prevent confident  identi fication of a rate-limiting mechanism from the  temperature  dependence of  the  growth rate of  the SiC depletion layer.  However,  linear-reaction-controlled  kinetics  is  unlikely  at  these  high  temperatures.  In  addition,  linear-diffusion-con trolled  depletion  layer  growth  kinetics  is  rejected  as  this  would imply a constant diffusion length for  the oxygen, and  this work  has  shown  that  the  borosilicate  glass  layer,  the  ZrO2 increasing  layer and the SiC depletion layer all grow with time,  the diffusion lengths  for O2(g) concluded that parabolic, gas-phase  inward or SiO(g)  outward (Fig. 11).  It  is  diffusion-limited growth of  the depletion layer best  explains  the observed results. As  in the case of  the transition from C  formation to SiC depletion,  it was again necessary to con sider  thermodynamics  at  the  local microstructural  ZrB2/ ZrO2/SiO2 junction to explain how O2(g) diffusion can be the rate-limiting mechanism.  V.  Conclusions  ZrB2-30 vol% SiC oxidation behavior has been characterized at temperatures of 1300°C to 1800°C. For temperatures below 1627°C, a two layer oxide is formed with ZrO2 + C below a 1627°C and temperatures of  borosilicate  glass  layer. For  above,  a  three  layer  depletion/oxidation morphology was  observed, with SiC depletion occurring beneath the ZrO2 and borosilicate layers. The growth of the SiC depletion layer is  best explained assuming parabolic gas diffusion limited active  oxidation of SiC to SiO(g). Consideration of local microstruc tural equilibrium was required to explain the C formation and SiC depletion behavior observed below and above 1627°C.  Acknowledgments  The authors would like to acknowledge Eric Neuman and Dr. William Fahrenholtz at Missouri University of Science and Technology for the ZrB2-30 vol% SiC material. Initial funding for this work was provided by The National Hyper sonic Science Center-Materials and Structures. This study was presented in part  at Materials Science and Technology, Montreal, Canada, October, 2013.  (a)  (b)  Fig. 15.  Fracture section of ZrB2-30 vol% SiC oxidized in flowing O2 using resistive heating system at 50 min, showing increased bubbling and ZrO2 growth at temperatures above the transition temperature.  (a)1600°C for 55 min (b) 1650°C for  Fig. 16. Diagram illustrating dissociation of SiO2(s,l) O2(g), at 1650°C allowing for continued SiC depletion.  to SiO(g) and  1682  Journal of the American Ceramic Society—Shugart and Opila  Vol. 98, No. 5  \\x0c', 'May 2015  SiC Depletion in ZrB2-SiC  1683  References  1A. Paul, D. D. Jayaseelan, S. Venugopal, E. Zapata-Sovas, J. G. P. 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},{
  "_id": 240,
  "PDF": "Significant improvement of the short-term high-temperature oxidation resistance of dense monolithic HfC-SiC ceramic nanocomposites upon incorporation of Ta.pdf",
  "Text": "['Corrosion Science 145 (2018) 191-198  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Significant improvement of the short-term high-temperature oxidation resistance of dense monolithic HfC/SiC ceramic nanocomposites upon incorporation of Ta  T  Qingbo Wen, Ralf Riedel, Emanuel  Ionescu⁎  Technische Universität Darmstadt,  Institut  für Materialwissenschaft, Otto-Bernd-Str. 3, Darmstadt, D-64287 Germany  A R T I C L E  I N F O  A B S T R A C T  Keywords:  Ceramic High-temperature corrosion Kinetic parameters Oxidation Tantalum oxide  1.  Introduction  The short-term oxidation (i.e., exposure time of up to 20 h) of dense monolithic (Hf,Ta)C/SiC-based ceramic nanocomposites at temperatures from 1200 to 1500 °C is presented and discussed. The oxidation behavior of HfC/SiC is similar to that of other UHTCs, indicating that it is mainly determined by the HfC phase. Ta incorporation into HfC leads to a strong decrease of the parabolic oxidation rates (3-4 orders of magnitude). Unlike HfC/SiC, Ta-containing nanocomposites form continuous multiphasic scales consisting of HfO2, Ta2O5 and Hf6Ta2O17, which are able to densify quickly during the oxidation process.  Ultrahigh-temperature ceramics (UHTCs)attracted much attention in the past due to their ultrahigh melting temperatures (≥ 3000 °C), high hardness (≥ 20 GPa), high Young’s modulus (≥ 300 GPa), excellent chemical stability and high electrical conductivity ( 104 S/cm) [1-4]. UHTCs are typically non-oxide ceramics such as carbides, nitrides and borides of early transition metal, e.g., HfC, HfN, ZrC, TaC, HfB2, ZrB2 [5,6]. However, UHTCs are prone to high-temperature oxidation, which seriously limits their practical applications [2,7,8]. Significant oxidation of UHTCs occurs in ambient atmosphere at temperatures starting with 400  800 °C and typically leads to porous oxide-based scales that are not able to protect from bulk oxidation [9-11]. In order to improve the environmental stability of UHTCs within the context of (ultra)high-temperature applications, UHTC-based compo(e.g., UHTC (nano)composites [12-16] and CMCs [12,13]) were sites developed. Typically, a secondary silica-forming phase is added to the UHTC, which is expected to protect the UHTC phases by forming a dense Si-based scale with low oxygen diffusivity at emperatures up to 1600 °C [16-20]. Thus, in UHTC-based (nano)composites, the Si-based protective scale (e.g., SiO2, borosilicate, metal silicates) is supposed to protect them at temperatures up to 1600 °C; whereas beyond this temperature the UHTC phase is expected to deliver oxide scales being able to quickly densify and thus providing environmental robustness at ultrahigh temperatures [13,21-25]. Moreover, the use of silica formers  as secondary phase leads to a reduction of the density in UHTC-based (nano)composites [26,27]. Whereas, appropriate volume fractions of the highly conductive UHTC phase can improve/determine the electrical conductivity of the UHTC-based (nano)composites [4,28]. One typical issue has been identified within the context of developing UHTCs-based composites with improved robustness in harsh environment and this relates to the mismatch in the oxidation kinetics of the UHTC phase and that of the silicon carbide (or in general the silica former phase). Thus, UHTC phases such as group IV metal carbides, nitrides or borides were shown to be highly sensitive to oxidation in the temperature range from e.g. 400-600 °C up to ca. 1600-1800 °C, due to the fact that they generate upon oxidation porous scales which are not able to protect from further oxidation. For instance, HfC was shown to be actively oxidized and form a porous hafnia scale at temperature starting from ca. 400 °C [9]. Only at very high temperatures (e.g., beyond 1600-1800 °C), the oxide scales are able to sinter quickly enough in order to provide reasonable oxidation protection [11,29]. Silicon carbide shows a continuous silica formation at temperatures typically beyond 1000-1200 °C and thus can protect the UHTC phase from oxidation. Unfortunately, in the intermediate temperature region from 400 to 600 °C to 1000 °C UHTCs have a fundamental issue in oxidation/corrosion environments, as the SiC phase is not able to protect the UHTC phase effectively at temperatures below 1000 °C. Within this context, there is still a need for materials and microstructure design solutions in order to overcome the mentioned oxidation/corrosion sensitivity of UHTC-base composites at intermediate temperatures.  ⁎ Corresponding author. E-mail address: ionescu@materials.tu-darmstadt.de (E.  Ionescu).  https://doi.org/10.1016/j.corsci.2018.10.005 Received 22 May 2018; Received in revised form 2 October 2018; Accepted 5 October 2018  Available online 09 October 2018 0010-938X/ © 2018 Published by Elsevier Ltd.  \\x0c', 'Q. Wen et al.  Recently, a series of novel, dense monolithic UHTC/SiC ceramic nanocomposites based on solid-solution (Hf1-xTax)C dispersed within a SiC matrix were prepared via pyrolysis of suitable polymeric singlesource precursors followed by spark plasma sintering (SPS) [3,14,15,17]. The prepared nanocomposites were shown to possess a rather unique microstructure as well as outstanding high-temperature stability [12-14,17] and exciting electrical/dielectric properties [3,28] In the present work, the oxidation behavior of the mentioned UHTC/SiC ceramic nanocomposites at temperatures up to 1500 °C is described in detail and discussed within the context of other UHTCbased materials and composites thereof reported in the literature. It is shown that alloying the metal carbide phase (i.e., HfC as for our study) with other transition metals (i.e., Ta as for our study) may be able to improve the oxidation behavior of the UHTC phases, especially in the intermediate temperature range.  2. Experimental procedure  2.1. Materials preparation  Dense monolithic TaC/SiC, (Hf0.2Ta0.8)C/SiC, (Hf0.7Ta0.3)C/SiC and HfC/SiC ceramic nanocomposites were prepared from the corresponding amorphous Six(Hf,Ta)yCz ceramic powders via spark plasma sintering (FCT HP D 25/1, FCT Systeme GmbH, Frankenblick, Germany for the Hf-containing monoliths and SPS-1050, SPS Syntex Inc., Kawasaki, Japan for monolithic TaC/SiC). The procedure involved a heating rate of 100 °C/min. to the target temperature of 2200 °C, holding time of 20 min under uniaxial pressure of 50 MPa and cooling with a rate of 220 °C. The ceramic samples were cut in coupons with dimensions of 4 × 3×2 mm3 and subsequently ground and polished with 1 μm polycrystalline diamond on felt cloth. The amorphous Six(Hf,Ta)yCz ceramic powders were synthesized via pyrolysis of Hfand Ta-containing polymeric precursors, as reported in [3,28]. The polymeric precursors were synthesized upon reacting a commercially available allylhydridopolycarbosilane (SMP10, Starfire System Inc, USA) with Hf(NEt2)4 and Ta(NMe2)5 (Merck, Germany) using a weight ratio amido complexes:SMP-10 of 30:70. [3,28] For the solid solution samples, i.e. (Hf0.2Ta0.8)C/SiC, (Hf0.7Ta0.3)C/SiC, the Hf:Ta molar ratio was set to 2:8 and 3:7, respectively.  2.2. Materials characterization  The as-prepared monolithic samples as well as the samples exposed to oxidation conditions were structurally characterized by elemental analysis, X-ray diffraction (XRD) and scanning electron microscopy (SEM). The carbon content in the samples was determined using a combustion analysis method on a LECO C-200 analyzer (LECO Corporation, St. Joseph, Michigan, USA). The nitrogen and oxygen contents were measured on a LECO TC-436 analyzer (LECO Corporation, St. Joseph, Michigan, USA). Hf, Ta, and Si elemental contents were determined using Inductively Coupled Plasma-Atomic Emission Spectrometry (ICP-AES, Mikroanalytisches Labor Pascher, Remagen, Germany). X-ray diffraction of powder samples was performed with a STADI P powder diffractometer (STOE & Cie GmbH, Kα1 Darmstadt, Germany, molybdenum radiation source, λ = 0.709300 Å). The crystalline phase composition of the monolithic samples was determined on a Bruker D8 system (Bruker Corporation, copper Kα1 Billerica, Massachusetts, United States, radiation source, λ = 1.541874 Å). SEM was done with a Philips XL30 FEG high-resolution scanning electron microscope (FEI Company, Hillsboro, Oregon, USA) coupled with energy-dispersive X-ray spectroscopy (EDAX, Mahwah, New Jersey, USA).  2.3.  Isothermal oxidation tests  The  oxidation experiments were performed in an alumina  tube  Corrosion Science 145 (2018) 191-198  furnace at temperatures of 1200-1500 °C in air environment. In order to prevent the introduction of impurities during the oxidation experiments, the samples were placed in SiC crucible. The monolithic samples were subjected to a thermal treatment (80 °C for 48 h) prior to the oxidation experiment. The oxidation experiments were performed as follows: the furnace was heated to the target temperature and held for 30 min; subsequently, the specimens were placed in the isothermal area of the furnace and their weight was measured after dwelling time of 1, 2, 3, 5, 10 and 20 h.  3. Results and discussion  Amorphous Si(Hf,Ta)C-based ceramic powders (i.e., SiHfC-1100 °C, SiTaC-1100 °C, Six(Hf0.7Ta0.3)yCz-1000 °C and Six(Hf0.2Ta0.8)yCz1000 °C) were obtained upon pyrolysis of the as-synthesized singlesource precursors in argon atmosphere at 1100 °C (for SixHfyCz and SixTayCz) or 1000 °C (Six(Hf0.7Ta0.3)yCz-1000 °C and Six(Hf0.2Ta0.8)yCz1000 °C) for 2 h [3,28]. The ceramic yield amounted ca. 80 wt.%. A detailed discussion on the structure of the used single-source precursors and their conversion into the amorphous Si(Hf,Ta)C ceramics has been recently published in [28]. After spark plasma sintering at 2200 °C, the X-ray amorphous ceramic powders (Fig. 1) undergo densification accompanied by partitioning and crystallization, thus allowing to obtain dense monolithic (Hf,Ta)C/SiC-based ceramic nanocomposites consisting of crystalline (Hf,Ta)C and β-SiC phases (Fig. 1) [3,28]. The XRD results (Fig. 1) indicate that in all prepared dense monolithic samples, nanoscaled (Hf,Ta)C phases were in situ generated within a β-SiC matrix during the SPS process. The average grain sizes, chemical composition (elemental analysis) and phase composition of the obtained monolithic samples are shown in Table 1. The estimated average β-SiC grains within grain sizes of both (Hf,Ta)C and all prepared monoliths are less than 100 nm. The open porosities of all monoliths are less than 1.5 vol.%, which indicates that they are dense and suitable for oxidation test. The phase composition (Table 1) shows that all monoliths possess comparable volume fractions of β-SiC, (Hf,Ta)C and segregated carbon. The dense monolithic Si(Hf,Ta)C-based nanocomposite samples were oxidized in air at 1200-1500 °C for 20 h. The photographs of the oxidized samples are shown in Figure S1 in the Supplementary Information. The HfC/SiC samples became gray and suffered from extensive cracking during the exposure to high temperatures and oxidative conditions; while the TaC/SiC samples showed no significant change in color or crack formation upon oxidation. Also, the other Ta Fig.  1. XRD patterns of the as-pyrolyzed amorphous Six(Hf,Ta)yCz-based ceramic powders (bottom) and of the as-sintered dense monolithic (Hf,Ta)C/ SiC-based ceramic nanocomposites (top) [3,28].  192  \\x0c', 'Q. Wen et al.  Table 1  Elemental composition, empirical nanocomposites.  formulae, phase composition and average grain sizes  for  (Hf,Ta)C and β-SiC in the studied (Hf,Ta)C/SiC monolithic ceramic  Corrosion Science 145 (2018) 191-198  Sample  Elemental contents [wt%]  Empirical  formulae  Phase composition [vol%]  Average grain size [nm]  TaC/SiC (Hf0.2Ta0.8)C/SiC (Hf0.7Ta0.3)C/SiC HfC/SiC  Si  52.0 51.7 50.9 49.1  Hf  —  3.4 13.3 21.9  Ta  18.0 14.3 6.0 —  C  26.5 27.5 27.0 25.1  N  0.2 < 0.1 < 0.1 0.4  O  3.9 3.1 2.3 3.5  SiTa0.05C1.2O0.13(N0.006) Si(Hf0.01Ta0.04)C1.24O0.1 Si(Hf0.04Ta0.02)C1.24O0.08 SiHf0.07C1.19O0.12(N0.02)  SiC  89.0 87.5 86.8 87.5  (Hf,Ta)C  4.9 4.9 5.8 7.0  C  6.1 7.7 7.4 5.5  (Hf,Ta)C  73.7 72.3 91.1 80.9  SiC  75.3 73.0 81.0 36.6  * The average grain sizes have been estimated by Rietveld refinement of  the XRD patterns using Full-Profile software [3,28].  Fig. 2.  Isothermal oxidation curves of the dense monoliths at temperatures from 1200 to 1500 °C: HfC/SiC (a), (Hf0.7Ta0.3)C/SiC (b), (Hf0.2Ta0.8)C/SiC (c) and TaC/ SiC (d).  containing samples (i.e., (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC) exhibited an improved oxidation behavior with respect to that of HfC/SiC; thus, no change in color and no obvious crack formation during the oxidation was observed. The oxidation curves for the Si(Hf,Ta)C-based monoliths oxidized at 1200 1500 °C 1200-1500 °C are shown in Fig. 2. The oxidation curves of (Hf0.7Ta0.3)C/SiC, (Hf0.2Ta0.8)C/SiC and TaC/SiC recorded at 1200 °C (Fig. 2) show that there is mass loss of the samples in the first 3 h of the process due to the oxidation of the segregated carbon present with fractions of ca. 5-8 vol% in the studied samples (Table 1). Subsequently, all samples show a parabolic oxidation behavior which is considered to rely on the formation and growth of a protective oxide scale. Thus, the apparent parabolic rate constant (Kp) can be calculated cf. , where, ∇W A/ is the surfacespecific mass change in mg/cm2, A is the specific surface area in cm2 and t is the oxidation time h [31]. It should be noted that the parabolic rates were determined from the mass change data in the oxidation time from 5 to 20 h. The data recorded at smaller soaking times (i.e., less  ∇(  W A  K  t*p  2  )  =  than 5 h) cannot be considered within the parabolic oxidation description, as mass loss was recorded due to the oxidation of segregated carbon. Consequently, those data were omitted for the calculation of the parabolic oxidation rates. The parabolic plots for the samples are shown in Fig. 3, which indicate that HfC/SiC exhibits apparent parabolic rate constants (Kp) of ca. 2-3 mg2/(cm4·h), which are similar to those reported for other Hfbased UHTCs (i.e., 2.56 and 2.1 mg2/(cm4·h) at 1500 °C for HfB2-20 vol % SiC [32] and HfC(N)/SiCN [15], respectively). Moreover, they are two to three orders of magnitude higher than those recorded at the same temperatures for CVD SiC [31,33-35] and sintered SiC [36] (Kp values in the order of 10−4 to 10-3 mg2/(cm4·h)). This indicates that the short-term oxidation kinetics (here, exposure time between 5 and 20 h) in HfC/SiC is mainly dictated by the HfC phase. All studied Ta-containing monoliths (including TaC/SiC, (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC) exhibit significantly higher oxidation resistance than that of HfC/SiC, independently of the oxidation temperature. The Kp values of all Ta-containing monoliths at 1300-1500 °C are comparable to those measured for the CVD SiC and  193  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 3. Parabolic plots of the surface specific mass change of the Si(Hf,Ta)C-based dense monoliths as function of the oxidation time: (a) HfC/SiC, (Hf0.7Ta0.3)C/SiC, (c) (Hf0.2Ta0.8)C/SiC, (d) TaC/SiC. The greyish area corresponds to the first 5 h of oxidation and was not considered for the calculation of the parabolic rate constants because of the mass loss related to segregated carbon (see Fig. 2). For the plots in (b), (c) and (d), the mass change in the curves at 1200 °C corresponds to the right y axes, whereas that in all other curves correspond to the left y axes.  sintered SiC upon oxidation at the same temperatures in dry oxygen. This is an intriguing result, as it is well known that both early transition metal carbides (e.g., TaC and HfC in the present case) exhibit poor oxidation resistance at the testing temperatures due to the formation of porous oxides [11,37,38]. This will be discussed in the following. The addition of Ta in the form of (Hf,Ta)C solid solution seems to significantly improve the oxidation behavior of the samples. The TaC/ SiC samples show consequently the lowest parabolic rates among the Ta-containing samples; the decrease of the Ta content (at the expenses of Hf) results in an increase of the Kp values. For instance, the Kp value of TaC/SiC measured upon oxidation at 1200 °C was 0.0123 mg2/(cm4 h); whereas the values for (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC at the same temperature were one order of magnitude higher, i.e., 0.28 and 0.38 mg2/(cm4 h), respectively. The highest Kp value at 1200 °C was recorded for HfC/SiC, i.e. 2.13 mg2/(cm4h). Additionally, an interesting trend has been observed in the case of all Ta-containing samples: The Kp values of the (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC monoliths at 1200 °C were found to be ca. 3 orders of magnitude higher than those obtained upon oxidation at higher temperatures, i.e. at 1300, 1400 and 1500 °C. Also, the Kp values of TaC/SiC at 1200-1500 °C decrease as the oxidation temperature increases, i.e. this is opposite to the behavior of pure SiC which shows an increase of the Kp values as the oxidation temperature increases. [33,39] This behavior was reported already in literature, e.g. for TaC, HfC or solid solution thereof, and was correlated to the improved sintering ability of the oxide scale as the oxidation temperature increases [29].  Kp1500 < Kp1400 < Kp1300 < Kp1200;  In order to rationalize the effect of Ta incorporation on the oxidation behavior of the studied (Hf,Ta)C/SiC ceramic nanocomposites and to assess the evolution of their surface during oxidation, the oxidized samples were investigated by means of XRD (Fig. 4). As the HfC/SiC samples were fully cracked after the oxidation process, their XRD patterns were recorded using powder X-ray diffraction (Fig. 4a). As shown in Fig. 4a, monoclinic HfO2 (m-HfO2) and small amounts of β-cristobalite (SiO2) are formed due to the oxidation of HfC and βSiC, respectively. Moreover, the formation of a minor amount of HfSiO4 is observed for the sample oxidized at 1500 °C, probably due to the reaction between SiO2 and HfO2 [8,40,41]. For all oxidized HfC/SiC samples less residual HfC can be observed. Thus, it may be concluded that the oxidation of the HfC phase in the studied temperature range is extensive and generates a porous m-HfO2 rather scale. Moreover, it seems to have the main contribution on the oxidation behavior of HfC/ rather high Kp values [10,11]. The SiC and consequently explains its oxidation of HfC nanoparticles induces significant surface cracking in the HfC/SiC samples (Fig. 5), which provides fast track pathways for oxygen penetrating into the monoliths and hence further accelerating the oxidation process [42]. This may be correlated with the temperature-dependent reversible m ↔ t phase transformation of HfO2, which is accompanied by large volume changes [42]. Despite the formation of SiO2 scale upon oxidation of the SiC matrix (which occurs parallely but is significantly slower), it is obvious that the oxide scale generated during the isothermal oxidation test does not effectively protect the HfC/SiC monoliths. The XRD patterns of the Ta-containing oxidized monoliths show the  194  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 4. XRD patterns of the surface of HfC/SiC (a), (Hf0.7Ta0.3)C/SiC (b), (Hf0.2Ta0.8)C/SiC (c) and TaC/SiC (d) after oxidation in air for 20 h at temperatures from 1200 to 1500 °C. The XRD patterns from (a) were recorded using a molybdenum Kα1 radiation source in a transmission mode (powder XRD); whereas the XRD patterns in (b), (c) and (d) were measured using a copper Kα1 radiation source in reflection mode]. In (b) and (c) also the XRD patterns of the as-prepared materials (i.e., prior to oxidation) are shown.  presence of several metal oxides (β-Ta2O5, Hf6Ta2O17 and m-HfO2) in addition to small amounts of β-cristobalite. This has been correlated with the different molar ratios of Hf to Ta within the (Hf,Ta)C phase as for (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC [43]. The oxidized the formation of Hf6Ta2O17 and β(Hf0.2Ta0.8)C/SiC samples exhibit Ta2O5 (though no m-HfO2) upon oxidation (Fig. 4c); whereas for the and m-HfO2 (Hf0.7Ta0.3)C/SiC sample, the formation of Hf6Ta2O17 (though no β-Ta2O5) has been shown upon oxidation (Fig. 4d). The formation of Hf6Ta2O17 as component of the scale during oxidation of (Hf1-xTax)C solid solution has been already reported in literature. It was shown, that Hf6Ta2O17 generation is favorable for solid solutions possessing large excess of Hf (typically, Hf6Ta2O17 is formed upon reaction of 1 mole of Ta2O5 with six equivalents of HfO2) [44,45]. This is in agreement with the results of the present study, as the amount of Hf6Ta2O17 formed upon oxidation of (Hf0.7Ta0.3)C/SiC at e.g. 1200 °C (Fig. 4b) was significantly higher than that formed via oxidation of (Hf0.2Ta0.8)C (Fig. 4c). Interestingly, all Ta-containing oxidized monoliths show the presence of high-intensity reflections of non-oxidized metal carbide phases, i.e., (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C as well as TaC, which seems to be generated during oxidation, as discussed below. The amount of the nonoxidized metal carbide phases is high especially in the samples exposed to high oxidation temperatures (i.e., 1400 and 1500 °C). This may be  clearly correlated to the relatively low Kp values of the Ta-containing monoliths oxidized at temperatures of 1300, 1400 and 1500 °C, which are comparable to those measured for pure SiC, despite the expected poor oxidation resistance of (Hf1-xTax)C. Whereas the samples oxidized at 1200 °C show significantly higher parabolic rates. This intriguing behavior which has already been mentioned above can be correlated to the phase composition and sintering ability of the oxidation scale in the Ta-containing nanocomposite samples, as elaborated in the following. The melting point of HfO2 and Ta2O5 are around 2800 °C and 1800 °C, respectively [37,46]. Consequently, it is expected that in the the β-Ta2O5 phase, which is generated during the case of TaC/SiC, oxidation at high temperatures, is able to densify quickly during the oxidation process, when the samples are exposed to temperature beyond 1300 °C; whereas a temperature of 1200 °C seems to be not sufficient to achieve a fast densification. Consequently, the oxidation of TaC/SiC at those high temperatures (i.e., > 1300 °C) may provide a continuous, dense scale which is able to protect the base material from further oxidation. Obviously, the Kp seems to correlate to the sintering kinetics of Ta2O5. However, the intrinsic behavior of the Ta2O5, which is generated upon oxidation of TaC, shows a good match with SiC concerning the coefficient of thermal expansion but undergoes a phase transformation at high temperatures (i.e. orthorhombic → tetragonal, i.e. β → α, at ca. 1360 °C) and thus suffers from big volume changes  195  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 5. SEM images of  the surface of the HfC/SiC, (Hf0.7Ta0.3)C/SiC, (Hf0.2Ta0.8)C/SiC and TaC/SiC monoliths oxidized at 1200, 1400 and 1500 °C in air for 20 h.  during the oxidation process [47]. There is no positive temperature dependence of the oxidation rate for HfC/SiC, as HfO2 is unable to densify in the temperature range of 1200-1500 °C. A case study related to the high temperature oxidation of HfC indicated that the oxidation kinetics of HfC is dominated by the oxygen diffusion through the porous HfO2 scale at temperatures lower than 1800 °C and is thus in agreement with our and others [46] observations; whereas at temperatures beyond 1800 °C the porous HfO2 scale starts to densify and the oxidation kinetics is determined by other mechanisms such as solid state (lattice) diffusion of oxygen through the scale [29]. Interestingly, (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C seem to behave similarly to TaC/SiC, indicating that the sintering of the scale during the oxidation process may have significant contribution to their oxidation kinetics. As previously reported in the literature, the melting point of HfO2 can be significantly reduced to temperatures below 2200 °C upon adding less than 10 mole% of Ta2O5 [48]. According to the molar ratio of Hf to Ta within the Ta-containing monoliths studied in the present work, the molar fraction of Ta2O5 within the oxide scale is larger than 10%. Thus, the melting point of the oxidation scale may be reduced  significantly with respect to the melting point of HfO2. Therefore, during the oxidation at 1300-1500 °C, the generated metal oxides (e.g., β-Ta2O5 and Hf6Ta2O17) can be sintered/densified effectively. Combined with the SiO2 being formed upon oxidation of SiC matrix, it is possible to form a continuous and protective oxide scale on the monoliths in a relatively short time. Thus, the intensities of the reflections of the remaining (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C phases after oxidation at 1400 and 1500 °C are stronger than those recorded in the samples oxidized at 1200 and 1300 °C (Fig. 4), thus being in agreement with the negative temperature their Kp dependence of (Fig. 3). At 1200 °C, the densification of the oxides in the scale is not sufficiently fast, thus significant surface cracking and porosity are observed (Fig. 5, Fig. 6). An interesting aspect related to the oxidation behavior of (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC concerns the formation of TaC during their oxidation, as mentioned above. Thus, it seems that the amount of TaC increases as the oxidation temperature of the nanocomposites increases. This clearly indicates that the formation of TaC should be considered as a consequence of exposure to oxidation conditions at high temperatures. Moreover, the amount of TaC increases  196  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 6. SEM micrographs of the cross sections of HfC/SiC, (Hf0.2Ta0.8)C/SiC and Tac/SiC after oxidation at 1300 °C for 20 h. It can be seen that HfC/SiC does not exhibit a continuous scale, thus having a high-thickness oxidized and porous area at the surface; whereas (Hf0.2Ta0.8)C/SiC and TaC/SiC clearly show the formation of a continuous scale upon oxidation.  from (Hf0.7Ta0.3)C/SiC to (Hf0.2Ta0.8)C/SiC, clearly showing a direct correlation between the content of Ta in the (HfxTa1-x)C phase of the nanocomposites and the amount of TaC generated upon oxidation. It has been reported in literature that the oxidation of (Hf1-xTax)C leads in a first step to the generation of Ta2O5 and HfO2. If no sufficient HfO2 is present in order to consume Ta2O5 upon formation of Hf6Ta2O17, excess Ta2O5 is able to react with HfC to deliver HfO2 and TaC [44,45]. Thus, the evolution of the (Hf1-xTax)C/SiC nanocomposites at hightemperatures in oxidative conditions may be described according to the phase composition of their scale by following processes:  30 ( Ta  Hf  0.2  )  C  0.8  51 O  2  +  →  40 ( Hf  0.7  Ta  )  C  0.3  63 O  2  +  →  11 Ta O  2  5  +  Hf Ta O  2  6  17  +  30 CO  4 Hf Ta O  2  6  17  +  2 Ta O  2  5  4 HfO  2  40 CO  +  +  3  Ta O  2  5  7  HfC  →  7  HfO  2  +  6  TaC  +  CO  +  2 SiC  +  3 O  2  →  2 SiO  2  2 CO  +  (1)  (2)  (3)  (4)  The process shown in Eq. (3), which is considered to be responsible for the generation of TaC during the high-temperature oxidation of (Hf1-xTax)C/SiC, requires the presence of HfC in contact with Ta2O5. However, no phase separation occurs in both (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C phases even after annealing of the nanocomposites at 1900 °C in Argon or during SPS at 2200 °C in vacuum [28]. Thus, it is considered that during the oxidation process Ta either selectively diffuses out of the (Hf1-xTax)C phase or it is oxidized much faster than Hf, thus HfC-rich nanoparticles are generated via depletion of Ta from (Hf1xTax)C, which can subsequently react with Ta2O5 to form TaC and HfO2. Similar phenomena were observed during the oxidation of some solid solutions and alloys, such as Ti and Cr depletion of HfC-based solid solutions [49], outward diffusion of Ti from Ti3SiC2 [50] as well as outward diffusion of Hf from Hf-containing Ni-based superalloy [51]. Based on the discussed oxidation of (Hf1-xTax)C/SiC, it is considered that a selective activation of the processes generating Hf6Ta2O17 and TaC along with the formation of silica upon SiC oxidation may be helpful to improve the oxidation resistance of the materials studied in the present work.  4. Conclusions and outlook  The isothermal oxidation behavior of dense monolithic (Hf,Ta)C/ SiC-based ceramic nanocomposites has been investigated at temperatures up to 1500 °C in dry air. The isothermal oxidation test shows that the Ta-containing monoliths [i.e., TaC/SiC, (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC] exhibit much lower apparent parabolic oxidation rates than those of the HfC/SiC monoliths. Following conclusions can be drawn from the present study:  (i) The incorporation of Ta into HfC leads ability of the oxide scale  to an enhanced sintering  197  (ii) During the oxidation of the (Hf1-xTax)C/SiC nanocomposites, the formation of Hf6Ta2O17 and TaC at the expenses of Ta2O5 and HfO2 has been observed. As Ta2O5 and HfO2 are known to undergo upon thermal annealing reversible phase transformation, their consumption is assumed to improve the stability of the oxide scale against cracking. Moreover, Hf6Ta2O17 exhibit relatively low oxygen diffusivity and does not undergo any phase transformation up to its melting point, thus being beneficial for the stability of the oxide scale. (iii) Based on the observed oxidation behavior of (Hf1-xTax)C/SiC-based nanocomposites, it is proposed to further improve their oxidation resistance via a short pre-oxidation step at temperatures sufficiently high to activate the formation of silica.  Acknowledgements  The authors thank Prof. Dr. Olivier Guillon (IEK-1, FZ Jülich, Germany) and Dr. Koji Morita (NIMS, Tsukuba, Japan) for support concerning the Spark Plasma Sintering experiments as well as Cong Zhou and Jean-Christophe Jaud (TU Darmstadt) for support with the elemental analyses and XRD measurements, respectively. Furthermore, support from the German Science Foundation (DFG, Bonn) as well as “Smart EU COST Action CM1302 (European network Inorganic Polymers”, SIPs) is acknowledged.  Appendix A. 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},{
  "_id": 241,
  "PDF": "Significant improvement of the short-term high-temperature oxidation resistance of dense monolithic HfCSiC ceramic nanocomposites upon incorporation of Ta.pdf",
  "Text": "['Corrosion Science 145 (2018) 191-198  Contents lists available at ScienceDirect  Corrosion Science  jou rna l homepage : www .e l sev ie r .com / loca te /co rsc i  Significant improvement of the short-term high-temperature oxidation resistance of dense monolithic HfC/SiC ceramic nanocomposites upon incorporation of Ta  T  Qingbo Wen, Ralf Riedel, Emanuel  Ionescu⁎  Technische Universität Darmstadt,  Institut  für Materialwissenschaft, Otto-Bernd-Str. 3, Darmstadt, D-64287 Germany  A R T I C L E  I N F O  A B S T R A C T  Keywords:  Ceramic High-temperature corrosion Kinetic parameters Oxidation Tantalum oxide  1.  Introduction  The short-term oxidation (i.e., exposure time of up to 20 h) of dense monolithic (Hf,Ta)C/SiC-based ceramic nanocomposites at temperatures from 1200 to 1500 °C is presented and discussed. The oxidation behavior of HfC/SiC is similar to that of other UHTCs, indicating that it is mainly determined by the HfC phase. Ta incorporation into HfC leads to a strong decrease of the parabolic oxidation rates (3-4 orders of magnitude). Unlike HfC/SiC, Ta-containing nanocomposites form continuous multiphasic scales consisting of HfO2, Ta2O5 and Hf6Ta2O17, which are able to densify quickly during the oxidation process.  Ultrahigh-temperature ceramics (UHTCs)attracted much attention in the past due to their ultrahigh melting temperatures (≥ 3000 °C), high hardness (≥ 20 GPa), high Young’s modulus (≥ 300 GPa), excellent chemical stability and high electrical conductivity ( 104 S/cm) [1-4]. UHTCs are typically non-oxide ceramics such as carbides, nitrides and borides of early transition metal, e.g., HfC, HfN, ZrC, TaC, HfB2, ZrB2 [5,6]. However, UHTCs are prone to high-temperature oxidation, which seriously limits their practical applications [2,7,8]. Significant oxidation of UHTCs occurs in ambient atmosphere at temperatures starting with 400  800 °C and typically leads to porous oxide-based scales that are not able to protect from bulk oxidation [9-11]. In order to improve the environmental stability of UHTCs within the context of (ultra)high-temperature applications, UHTC-based compo(e.g., UHTC (nano)composites [12-16] and CMCs [12,13]) were sites developed. Typically, a secondary silica-forming phase is added to the UHTC, which is expected to protect the UHTC phases by forming a dense Si-based scale with low oxygen diffusivity at emperatures up to 1600 °C [16-20]. Thus, in UHTC-based (nano)composites, the Si-based protective scale (e.g., SiO2, borosilicate, metal silicates) is supposed to protect them at temperatures up to 1600 °C; whereas beyond this temperature the UHTC phase is expected to deliver oxide scales being able to quickly densify and thus providing environmental robustness at ultrahigh temperatures [13,21-25]. Moreover, the use of silica formers  as secondary phase leads to a reduction of the density in UHTC-based (nano)composites [26,27]. Whereas, appropriate volume fractions of the highly conductive UHTC phase can improve/determine the electrical conductivity of the UHTC-based (nano)composites [4,28]. One typical issue has been identified within the context of developing UHTCs-based composites with improved robustness in harsh environment and this relates to the mismatch in the oxidation kinetics of the UHTC phase and that of the silicon carbide (or in general the silica former phase). Thus, UHTC phases such as group IV metal carbides, nitrides or borides were shown to be highly sensitive to oxidation in the temperature range from e.g. 400-600 °C up to ca. 1600-1800 °C, due to the fact that they generate upon oxidation porous scales which are not able to protect from further oxidation. For instance, HfC was shown to be actively oxidized and form a porous hafnia scale at temperature starting from ca. 400 °C [9]. Only at very high temperatures (e.g., beyond 1600-1800 °C), the oxide scales are able to sinter quickly enough in order to provide reasonable oxidation protection [11,29]. Silicon carbide shows a continuous silica formation at temperatures typically beyond 1000-1200 °C and thus can protect the UHTC phase from oxidation. Unfortunately, in the intermediate temperature region from 400 to 600 °C to 1000 °C UHTCs have a fundamental issue in oxidation/corrosion environments, as the SiC phase is not able to protect the UHTC phase effectively at temperatures below 1000 °C. Within this context, there is still a need for materials and microstructure design solutions in order to overcome the mentioned oxidation/corrosion sensitivity of UHTC-base composites at intermediate temperatures.  ⁎ Corresponding author. E-mail address: ionescu@materials.tu-darmstadt.de (E.  Ionescu).  https://doi.org/10.1016/j.corsci.2018.10.005 Received 22 May 2018; Received in revised form 2 October 2018; Accepted 5 October 2018  Available online 09 October 2018 0010-938X/ © 2018 Published by Elsevier Ltd.  \\x0c', 'Q. Wen et al.  Recently, a series of novel, dense monolithic UHTC/SiC ceramic nanocomposites based on solid-solution (Hf1-xTax)C dispersed within a SiC matrix were prepared via pyrolysis of suitable polymeric singlesource precursors followed by spark plasma sintering (SPS) [3,14,15,17]. The prepared nanocomposites were shown to possess a rather unique microstructure as well as outstanding high-temperature stability [12-14,17] and exciting electrical/dielectric properties [3,28] In the present work, the oxidation behavior of the mentioned UHTC/SiC ceramic nanocomposites at temperatures up to 1500 °C is described in detail and discussed within the context of other UHTCbased materials and composites thereof reported in the literature. It is shown that alloying the metal carbide phase (i.e., HfC as for our study) with other transition metals (i.e., Ta as for our study) may be able to improve the oxidation behavior of the UHTC phases, especially in the intermediate temperature range.  2. Experimental procedure  2.1. Materials preparation  Dense monolithic TaC/SiC, (Hf0.2Ta0.8)C/SiC, (Hf0.7Ta0.3)C/SiC and HfC/SiC ceramic nanocomposites were prepared from the corresponding amorphous Six(Hf,Ta)yCz ceramic powders via spark plasma sintering (FCT HP D 25/1, FCT Systeme GmbH, Frankenblick, Germany for the Hf-containing monoliths and SPS-1050, SPS Syntex Inc., Kawasaki, Japan for monolithic TaC/SiC). The procedure involved a heating rate of 100 °C/min. to the target temperature of 2200 °C, holding time of 20 min under uniaxial pressure of 50 MPa and cooling with a rate of 220 °C. The ceramic samples were cut in coupons with dimensions of 4 × 3×2 mm3 and subsequently ground and polished with 1 μm polycrystalline diamond on felt cloth. The amorphous Six(Hf,Ta)yCz ceramic powders were synthesized via pyrolysis of Hfand Ta-containing polymeric precursors, as reported in [3,28]. The polymeric precursors were synthesized upon reacting a commercially available allylhydridopolycarbosilane (SMP10, Starfire System Inc, USA) with Hf(NEt2)4 and Ta(NMe2)5 (Merck, Germany) using a weight ratio amido complexes:SMP-10 of 30:70. [3,28] For the solid solution samples, i.e. (Hf0.2Ta0.8)C/SiC, (Hf0.7Ta0.3)C/SiC, the Hf:Ta molar ratio was set to 2:8 and 3:7, respectively.  2.2. Materials characterization  The as-prepared monolithic samples as well as the samples exposed to oxidation conditions were structurally characterized by elemental analysis, X-ray diffraction (XRD) and scanning electron microscopy (SEM). The carbon content in the samples was determined using a combustion analysis method on a LECO C-200 analyzer (LECO Corporation, St. Joseph, Michigan, USA). The nitrogen and oxygen contents were measured on a LECO TC-436 analyzer (LECO Corporation, St. Joseph, Michigan, USA). Hf, Ta, and Si elemental contents were determined using Inductively Coupled Plasma-Atomic Emission Spectrometry (ICP-AES, Mikroanalytisches Labor Pascher, Remagen, Germany). X-ray diffraction of powder samples was performed with a STADI P powder diffractometer (STOE & Cie GmbH, Kα1 Darmstadt, Germany, molybdenum radiation source, λ = 0.709300 Å). The crystalline phase composition of the monolithic samples was determined on a Bruker D8 system (Bruker Corporation, copper Kα1 Billerica, Massachusetts, United States, radiation source, λ = 1.541874 Å). SEM was done with a Philips XL30 FEG high-resolution scanning electron microscope (FEI Company, Hillsboro, Oregon, USA) coupled with energy-dispersive X-ray spectroscopy (EDAX, Mahwah, New Jersey, USA).  2.3.  Isothermal oxidation tests  The  oxidation experiments were performed in an alumina  tube  Corrosion Science 145 (2018) 191-198  furnace at temperatures of 1200-1500 °C in air environment. In order to prevent the introduction of impurities during the oxidation experiments, the samples were placed in SiC crucible. The monolithic samples were subjected to a thermal treatment (80 °C for 48 h) prior to the oxidation experiment. The oxidation experiments were performed as follows: the furnace was heated to the target temperature and held for 30 min; subsequently, the specimens were placed in the isothermal area of the furnace and their weight was measured after dwelling time of 1, 2, 3, 5, 10 and 20 h.  3. Results and discussion  Amorphous Si(Hf,Ta)C-based ceramic powders (i.e., SiHfC-1100 °C, SiTaC-1100 °C, Six(Hf0.7Ta0.3)yCz-1000 °C and Six(Hf0.2Ta0.8)yCz1000 °C) were obtained upon pyrolysis of the as-synthesized singlesource precursors in argon atmosphere at 1100 °C (for SixHfyCz and SixTayCz) or 1000 °C (Six(Hf0.7Ta0.3)yCz-1000 °C and Six(Hf0.2Ta0.8)yCz1000 °C) for 2 h [3,28]. The ceramic yield amounted ca. 80 wt.%. A detailed discussion on the structure of the used single-source precursors and their conversion into the amorphous Si(Hf,Ta)C ceramics has been recently published in [28]. After spark plasma sintering at 2200 °C, the X-ray amorphous ceramic powders (Fig. 1) undergo densification accompanied by partitioning and crystallization, thus allowing to obtain dense monolithic (Hf,Ta)C/SiC-based ceramic nanocomposites consisting of crystalline (Hf,Ta)C and β-SiC phases (Fig. 1) [3,28]. The XRD results (Fig. 1) indicate that in all prepared dense monolithic samples, nanoscaled (Hf,Ta)C phases were in situ generated within a β-SiC matrix during the SPS process. The average grain sizes, chemical composition (elemental analysis) and phase composition of the obtained monolithic samples are shown in Table 1. The estimated average β-SiC grains within grain sizes of both (Hf,Ta)C and all prepared monoliths are less than 100 nm. The open porosities of all monoliths are less than 1.5 vol.%, which indicates that they are dense and suitable for oxidation test. The phase composition (Table 1) shows that all monoliths possess comparable volume fractions of β-SiC, (Hf,Ta)C and segregated carbon. The dense monolithic Si(Hf,Ta)C-based nanocomposite samples were oxidized in air at 1200-1500 °C for 20 h. The photographs of the oxidized samples are shown in Figure S1 in the Supplementary Information. The HfC/SiC samples became gray and suffered from extensive cracking during the exposure to high temperatures and oxidative conditions; while the TaC/SiC samples showed no significant change in color or crack formation upon oxidation. Also, the other Ta Fig.  1. XRD patterns of the as-pyrolyzed amorphous Six(Hf,Ta)yCz-based ceramic powders (bottom) and of the as-sintered dense monolithic (Hf,Ta)C/ SiC-based ceramic nanocomposites (top) [3,28].  192  \\x0c', 'Q. Wen et al.  Table 1  Elemental composition, empirical nanocomposites.  formulae, phase composition and average grain sizes  for  (Hf,Ta)C and β-SiC in the studied (Hf,Ta)C/SiC monolithic ceramic  Corrosion Science 145 (2018) 191-198  Sample  Elemental contents [wt%]  Empirical  formulae  Phase composition [vol%]  Average grain size [nm]  TaC/SiC (Hf0.2Ta0.8)C/SiC (Hf0.7Ta0.3)C/SiC HfC/SiC  Si  52.0 51.7 50.9 49.1  Hf  —  3.4 13.3 21.9  Ta  18.0 14.3 6.0 —  C  26.5 27.5 27.0 25.1  N  0.2 < 0.1 < 0.1 0.4  O  3.9 3.1 2.3 3.5  SiTa0.05C1.2O0.13(N0.006) Si(Hf0.01Ta0.04)C1.24O0.1 Si(Hf0.04Ta0.02)C1.24O0.08 SiHf0.07C1.19O0.12(N0.02)  SiC  89.0 87.5 86.8 87.5  (Hf,Ta)C  4.9 4.9 5.8 7.0  C  6.1 7.7 7.4 5.5  (Hf,Ta)C  73.7 72.3 91.1 80.9  SiC  75.3 73.0 81.0 36.6  * The average grain sizes have been estimated by Rietveld refinement of  the XRD patterns using Full-Profile software [3,28].  Fig. 2.  Isothermal oxidation curves of the dense monoliths at temperatures from 1200 to 1500 °C: HfC/SiC (a), (Hf0.7Ta0.3)C/SiC (b), (Hf0.2Ta0.8)C/SiC (c) and TaC/ SiC (d).  containing samples (i.e., (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC) exhibited an improved oxidation behavior with respect to that of HfC/SiC; thus, no change in color and no obvious crack formation during the oxidation was observed. The oxidation curves for the Si(Hf,Ta)C-based monoliths oxidized at 1200 1500 °C 1200-1500 °C are shown in Fig. 2. The oxidation curves of (Hf0.7Ta0.3)C/SiC, (Hf0.2Ta0.8)C/SiC and TaC/SiC recorded at 1200 °C (Fig. 2) show that there is mass loss of the samples in the first 3 h of the process due to the oxidation of the segregated carbon present with fractions of ca. 5-8 vol% in the studied samples (Table 1). Subsequently, all samples show a parabolic oxidation behavior which is considered to rely on the formation and growth of a protective oxide scale. Thus, the apparent parabolic rate constant (Kp) can be calculated cf. , where, ∇W A/ is the surfacespecific mass change in mg/cm2, A is the specific surface area in cm2 and t is the oxidation time h [31]. It should be noted that the parabolic rates were determined from the mass change data in the oxidation time from 5 to 20 h. The data recorded at smaller soaking times (i.e., less  ∇(  W A  K  t*p  2  )  =  than 5 h) cannot be considered within the parabolic oxidation description, as mass loss was recorded due to the oxidation of segregated carbon. Consequently, those data were omitted for the calculation of the parabolic oxidation rates. The parabolic plots for the samples are shown in Fig. 3, which indicate that HfC/SiC exhibits apparent parabolic rate constants (Kp) of ca. 2-3 mg2/(cm4·h), which are similar to those reported for other Hfbased UHTCs (i.e., 2.56 and 2.1 mg2/(cm4·h) at 1500 °C for HfB2-20 vol % SiC [32] and HfC(N)/SiCN [15], respectively). Moreover, they are two to three orders of magnitude higher than those recorded at the same temperatures for CVD SiC [31,33-35] and sintered SiC [36] (Kp values in the order of 10−4 to 10-3 mg2/(cm4·h)). This indicates that the short-term oxidation kinetics (here, exposure time between 5 and 20 h) in HfC/SiC is mainly dictated by the HfC phase. All studied Ta-containing monoliths (including TaC/SiC, (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC) exhibit significantly higher oxidation resistance than that of HfC/SiC, independently of the oxidation temperature. The Kp values of all Ta-containing monoliths at 1300-1500 °C are comparable to those measured for the CVD SiC and  193  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 3. Parabolic plots of the surface specific mass change of the Si(Hf,Ta)C-based dense monoliths as function of the oxidation time: (a) HfC/SiC, (Hf0.7Ta0.3)C/SiC, (c) (Hf0.2Ta0.8)C/SiC, (d) TaC/SiC. The greyish area corresponds to the first 5 h of oxidation and was not considered for the calculation of the parabolic rate constants because of the mass loss related to segregated carbon (see Fig. 2). For the plots in (b), (c) and (d), the mass change in the curves at 1200 °C corresponds to the right y axes, whereas that in all other curves correspond to the left y axes.  sintered SiC upon oxidation at the same temperatures in dry oxygen. This is an intriguing result, as it is well known that both early transition metal carbides (e.g., TaC and HfC in the present case) exhibit poor oxidation resistance at the testing temperatures due to the formation of porous oxides [11,37,38]. This will be discussed in the following. The addition of Ta in the form of (Hf,Ta)C solid solution seems to significantly improve the oxidation behavior of the samples. The TaC/ SiC samples show consequently the lowest parabolic rates among the Ta-containing samples; the decrease of the Ta content (at the expenses of Hf) results in an increase of the Kp values. For instance, the Kp value of TaC/SiC measured upon oxidation at 1200 °C was 0.0123 mg2/(cm4 h); whereas the values for (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC at the same temperature were one order of magnitude higher, i.e., 0.28 and 0.38 mg2/(cm4 h), respectively. The highest Kp value at 1200 °C was recorded for HfC/SiC, i.e. 2.13 mg2/(cm4h). Additionally, an interesting trend has been observed in the case of all Ta-containing samples: The Kp values of the (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC monoliths at 1200 °C were found to be ca. 3 orders of magnitude higher than those obtained upon oxidation at higher temperatures, i.e. at 1300, 1400 and 1500 °C. Also, the Kp values of TaC/SiC at 1200-1500 °C decrease as the oxidation temperature increases, i.e. this is opposite to the behavior of pure SiC which shows an increase of the Kp values as the oxidation temperature increases. [33,39] This behavior was reported already in literature, e.g. for TaC, HfC or solid solution thereof, and was correlated to the improved sintering ability of the oxide scale as the oxidation temperature increases [29].  Kp1500 < Kp1400 < Kp1300 < Kp1200;  In order to rationalize the effect of Ta incorporation on the oxidation behavior of the studied (Hf,Ta)C/SiC ceramic nanocomposites and to assess the evolution of their surface during oxidation, the oxidized samples were investigated by means of XRD (Fig. 4). As the HfC/SiC samples were fully cracked after the oxidation process, their XRD patterns were recorded using powder X-ray diffraction (Fig. 4a). As shown in Fig. 4a, monoclinic HfO2 (m-HfO2) and small amounts of β-cristobalite (SiO2) are formed due to the oxidation of HfC and βSiC, respectively. Moreover, the formation of a minor amount of HfSiO4 is observed for the sample oxidized at 1500 °C, probably due to the reaction between SiO2 and HfO2 [8,40,41]. For all oxidized HfC/SiC samples less residual HfC can be observed. Thus, it may be concluded that the oxidation of the HfC phase in the studied temperature range is extensive and generates a porous m-HfO2 rather scale. Moreover, it seems to have the main contribution on the oxidation behavior of HfC/ rather high Kp values [10,11]. The SiC and consequently explains its oxidation of HfC nanoparticles induces significant surface cracking in the HfC/SiC samples (Fig. 5), which provides fast track pathways for oxygen penetrating into the monoliths and hence further accelerating the oxidation process [42]. This may be correlated with the temperature-dependent reversible m ↔ t phase transformation of HfO2, which is accompanied by large volume changes [42]. Despite the formation of SiO2 scale upon oxidation of the SiC matrix (which occurs parallely but is significantly slower), it is obvious that the oxide scale generated during the isothermal oxidation test does not effectively protect the HfC/SiC monoliths. The XRD patterns of the Ta-containing oxidized monoliths show the  194  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 4. XRD patterns of the surface of HfC/SiC (a), (Hf0.7Ta0.3)C/SiC (b), (Hf0.2Ta0.8)C/SiC (c) and TaC/SiC (d) after oxidation in air for 20 h at temperatures from 1200 to 1500 °C. The XRD patterns from (a) were recorded using a molybdenum Kα1 radiation source in a transmission mode (powder XRD); whereas the XRD patterns in (b), (c) and (d) were measured using a copper Kα1 radiation source in reflection mode]. In (b) and (c) also the XRD patterns of the as-prepared materials (i.e., prior to oxidation) are shown.  presence of several metal oxides (β-Ta2O5, Hf6Ta2O17 and m-HfO2) in addition to small amounts of β-cristobalite. This has been correlated with the different molar ratios of Hf to Ta within the (Hf,Ta)C phase as for (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC [43]. The oxidized the formation of Hf6Ta2O17 and β(Hf0.2Ta0.8)C/SiC samples exhibit Ta2O5 (though no m-HfO2) upon oxidation (Fig. 4c); whereas for the and m-HfO2 (Hf0.7Ta0.3)C/SiC sample, the formation of Hf6Ta2O17 (though no β-Ta2O5) has been shown upon oxidation (Fig. 4d). The formation of Hf6Ta2O17 as component of the scale during oxidation of (Hf1-xTax)C solid solution has been already reported in literature. It was shown, that Hf6Ta2O17 generation is favorable for solid solutions possessing large excess of Hf (typically, Hf6Ta2O17 is formed upon reaction of 1 mole of Ta2O5 with six equivalents of HfO2) [44,45]. This is in agreement with the results of the present study, as the amount of Hf6Ta2O17 formed upon oxidation of (Hf0.7Ta0.3)C/SiC at e.g. 1200 °C (Fig. 4b) was significantly higher than that formed via oxidation of (Hf0.2Ta0.8)C (Fig. 4c). Interestingly, all Ta-containing oxidized monoliths show the presence of high-intensity reflections of non-oxidized metal carbide phases, i.e., (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C as well as TaC, which seems to be generated during oxidation, as discussed below. The amount of the nonoxidized metal carbide phases is high especially in the samples exposed to high oxidation temperatures (i.e., 1400 and 1500 °C). This may be  clearly correlated to the relatively low Kp values of the Ta-containing monoliths oxidized at temperatures of 1300, 1400 and 1500 °C, which are comparable to those measured for pure SiC, despite the expected poor oxidation resistance of (Hf1-xTax)C. Whereas the samples oxidized at 1200 °C show significantly higher parabolic rates. This intriguing behavior which has already been mentioned above can be correlated to the phase composition and sintering ability of the oxidation scale in the Ta-containing nanocomposite samples, as elaborated in the following. The melting point of HfO2 and Ta2O5 are around 2800 °C and 1800 °C, respectively [37,46]. Consequently, it is expected that in the the β-Ta2O5 phase, which is generated during the case of TaC/SiC, oxidation at high temperatures, is able to densify quickly during the oxidation process, when the samples are exposed to temperature beyond 1300 °C; whereas a temperature of 1200 °C seems to be not sufficient to achieve a fast densification. Consequently, the oxidation of TaC/SiC at those high temperatures (i.e., > 1300 °C) may provide a continuous, dense scale which is able to protect the base material from further oxidation. Obviously, the Kp seems to correlate to the sintering kinetics of Ta2O5. However, the intrinsic behavior of the Ta2O5, which is generated upon oxidation of TaC, shows a good match with SiC concerning the coefficient of thermal expansion but undergoes a phase transformation at high temperatures (i.e. orthorhombic → tetragonal, i.e. β → α, at ca. 1360 °C) and thus suffers from big volume changes  195  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 5. SEM images of  the surface of the HfC/SiC, (Hf0.7Ta0.3)C/SiC, (Hf0.2Ta0.8)C/SiC and TaC/SiC monoliths oxidized at 1200, 1400 and 1500 °C in air for 20 h.  during the oxidation process [47]. There is no positive temperature dependence of the oxidation rate for HfC/SiC, as HfO2 is unable to densify in the temperature range of 1200-1500 °C. A case study related to the high temperature oxidation of HfC indicated that the oxidation kinetics of HfC is dominated by the oxygen diffusion through the porous HfO2 scale at temperatures lower than 1800 °C and is thus in agreement with our and others [46] observations; whereas at temperatures beyond 1800 °C the porous HfO2 scale starts to densify and the oxidation kinetics is determined by other mechanisms such as solid state (lattice) diffusion of oxygen through the scale [29]. Interestingly, (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C seem to behave similarly to TaC/SiC, indicating that the sintering of the scale during the oxidation process may have significant contribution to their oxidation kinetics. As previously reported in the literature, the melting point of HfO2 can be significantly reduced to temperatures below 2200 °C upon adding less than 10 mole% of Ta2O5 [48]. According to the molar ratio of Hf to Ta within the Ta-containing monoliths studied in the present work, the molar fraction of Ta2O5 within the oxide scale is larger than 10%. Thus, the melting point of the oxidation scale may be reduced  significantly with respect to the melting point of HfO2. Therefore, during the oxidation at 1300-1500 °C, the generated metal oxides (e.g., β-Ta2O5 and Hf6Ta2O17) can be sintered/densified effectively. Combined with the SiO2 being formed upon oxidation of SiC matrix, it is possible to form a continuous and protective oxide scale on the monoliths in a relatively short time. Thus, the intensities of the reflections of the remaining (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C phases after oxidation at 1400 and 1500 °C are stronger than those recorded in the samples oxidized at 1200 and 1300 °C (Fig. 4), thus being in agreement with the negative temperature their Kp dependence of (Fig. 3). At 1200 °C, the densification of the oxides in the scale is not sufficiently fast, thus significant surface cracking and porosity are observed (Fig. 5, Fig. 6). An interesting aspect related to the oxidation behavior of (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC concerns the formation of TaC during their oxidation, as mentioned above. Thus, it seems that the amount of TaC increases as the oxidation temperature of the nanocomposites increases. This clearly indicates that the formation of TaC should be considered as a consequence of exposure to oxidation conditions at high temperatures. Moreover, the amount of TaC increases  196  \\x0c', 'Q. Wen et al.  Corrosion Science 145 (2018) 191-198  Fig. 6. SEM micrographs of the cross sections of HfC/SiC, (Hf0.2Ta0.8)C/SiC and Tac/SiC after oxidation at 1300 °C for 20 h. It can be seen that HfC/SiC does not exhibit a continuous scale, thus having a high-thickness oxidized and porous area at the surface; whereas (Hf0.2Ta0.8)C/SiC and TaC/SiC clearly show the formation of a continuous scale upon oxidation.  from (Hf0.7Ta0.3)C/SiC to (Hf0.2Ta0.8)C/SiC, clearly showing a direct correlation between the content of Ta in the (HfxTa1-x)C phase of the nanocomposites and the amount of TaC generated upon oxidation. It has been reported in literature that the oxidation of (Hf1-xTax)C leads in a first step to the generation of Ta2O5 and HfO2. If no sufficient HfO2 is present in order to consume Ta2O5 upon formation of Hf6Ta2O17, excess Ta2O5 is able to react with HfC to deliver HfO2 and TaC [44,45]. Thus, the evolution of the (Hf1-xTax)C/SiC nanocomposites at hightemperatures in oxidative conditions may be described according to the phase composition of their scale by following processes:  30 ( Ta  Hf  0.2  )  C  0.8  51 O  2  +  →  40 ( Hf  0.7  Ta  )  C  0.3  63 O  2  +  →  11 Ta O  2  5  +  Hf Ta O  2  6  17  +  30 CO  4 Hf Ta O  2  6  17  +  2 Ta O  2  5  4 HfO  2  40 CO  +  +  3  Ta O  2  5  7  HfC  →  7  HfO  2  +  6  TaC  +  CO  +  2 SiC  +  3 O  2  →  2 SiO  2  2 CO  +  (1)  (2)  (3)  (4)  The process shown in Eq. (3), which is considered to be responsible for the generation of TaC during the high-temperature oxidation of (Hf1-xTax)C/SiC, requires the presence of HfC in contact with Ta2O5. However, no phase separation occurs in both (Hf0.2Ta0.8)C and (Hf0.7Ta0.3)C phases even after annealing of the nanocomposites at 1900 °C in Argon or during SPS at 2200 °C in vacuum [28]. Thus, it is considered that during the oxidation process Ta either selectively diffuses out of the (Hf1-xTax)C phase or it is oxidized much faster than Hf, thus HfC-rich nanoparticles are generated via depletion of Ta from (Hf1xTax)C, which can subsequently react with Ta2O5 to form TaC and HfO2. Similar phenomena were observed during the oxidation of some solid solutions and alloys, such as Ti and Cr depletion of HfC-based solid solutions [49], outward diffusion of Ti from Ti3SiC2 [50] as well as outward diffusion of Hf from Hf-containing Ni-based superalloy [51]. Based on the discussed oxidation of (Hf1-xTax)C/SiC, it is considered that a selective activation of the processes generating Hf6Ta2O17 and TaC along with the formation of silica upon SiC oxidation may be helpful to improve the oxidation resistance of the materials studied in the present work.  4. Conclusions and outlook  The isothermal oxidation behavior of dense monolithic (Hf,Ta)C/ SiC-based ceramic nanocomposites has been investigated at temperatures up to 1500 °C in dry air. The isothermal oxidation test shows that the Ta-containing monoliths [i.e., TaC/SiC, (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC] exhibit much lower apparent parabolic oxidation rates than those of the HfC/SiC monoliths. Following conclusions can be drawn from the present study:  (i) The incorporation of Ta into HfC leads ability of the oxide scale  to an enhanced sintering  197  (ii) During the oxidation of the (Hf1-xTax)C/SiC nanocomposites, the formation of Hf6Ta2O17 and TaC at the expenses of Ta2O5 and HfO2 has been observed. As Ta2O5 and HfO2 are known to undergo upon thermal annealing reversible phase transformation, their consumption is assumed to improve the stability of the oxide scale against cracking. Moreover, Hf6Ta2O17 exhibit relatively low oxygen diffusivity and does not undergo any phase transformation up to its melting point, thus being beneficial for the stability of the oxide scale. (iii) Based on the observed oxidation behavior of (Hf1-xTax)C/SiC-based nanocomposites, it is proposed to further improve their oxidation resistance via a short pre-oxidation step at temperatures sufficiently high to activate the formation of silica.  Acknowledgements  The authors thank Prof. Dr. Olivier Guillon (IEK-1, FZ Jülich, Germany) and Dr. Koji Morita (NIMS, Tsukuba, Japan) for support concerning the Spark Plasma Sintering experiments as well as Cong Zhou and Jean-Christophe Jaud (TU Darmstadt) for support with the elemental analyses and XRD measurements, respectively. Furthermore, support from the German Science Foundation (DFG, Bonn) as well as “Smart EU COST Action CM1302 (European network Inorganic Polymers”, SIPs) is acknowledged.  Appendix A. 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  "_id": 242,
  "PDF": "Sintering and Mechanical Properties of ZrB2–TaSi2 and HfB2–TaSi2 Ceramic Composites.pdf",
  "Text": "['Sintering and Mechanical Properties of ZrB2-TaSi2 and HfB2-TaSi2  Ceramic Composites  Diletta Sciti,  w  Laura Silvestroni, Giancarlo Celotti, Cesare Melandri, and Stefano Guicciardi  CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy  Fully dense fine-grained ZrB2and HfB2-based composites contaning 15 vol% TaSi2 were produced by hot pressing at 18501- 19001C. Gas formation and mass loss, which occurred during sintering in  both  systems, were  in  agreement with  thermo dynamic predictions. In both composites, the presence of a solid  solution formed by the diffusion of  tantalum into the boride  matrix was observed. The HfB2-based composite was harder (22 GPa), stiffer (528 GPa), and tougher (4.1M Pa . m1/2) than the ZrB2-based composite. Although the room-temperature flexural strength of the ZrB2-based composite (830 MPa) was higher than that of the HfB2-based composite (700 MPa), the opposite was true at 12001 and 15001C. Contrary to the significant decrease observed for the ZrB2-based materials at the HfB2 composite retained B86% of elevated temperature, its room temperature strength up to 15001C (B600 MPa).  strength  1.  Introduction  Z IRCONIUM diboride and hafnium diboride belong to the class of ceramic materials known ultra-high-temperature-ceramics (UHTCs). These materials are of particular interest because  of  the unique combination of properties  they possess  such as  high refractoriness, high electrical  and thermal  conductivity,  chemical inertness against molten metals, or nonbasic and good oxidation resistance.1-10 Applications  slags,  that  take ad vantage of these properties include refractory linings, electrodes,  microelectronics,  and cutting  tools. Potential  applications of  zirconium and hafnium diboride also include aerospace manu facturing, for example, the sharp leading edge parts on hypersonic atmospheric reentry vehicles.1-10  Much effort has been devoted in recent years to improve the  densification and microstructure of these ultra refractory compounds.10-17 Previous studies have demonstrated that the addisuch as MoSi2 17-22 has positive effects on the densification and mechanical properties of borides, even at temperatures as high as 15001C. In addition, silicides, as silica-form tion of  silicides  ing  compounds,  can  offer  significant  improvement  to  the  oxidation behavior.  In this work the effect of TaSi2 addition on densification and properties of the borides is studied. So far, this silicide has been  employed for thin film applications in the field of electronics. To  the best of the authors’ knowledge, few data are available in the  literature on the properties of polycrystalline TaSi2 ceramics. TaSi2 has a melting point of 22001C, bulk density of 9.14 g/cm3, hardness  of 15.6 GPa, and is electrically conductive (electrical resistivity 50- 55 mO \\x01 cm).23 The elastic properties of TaSi2 single crystals were investigated by Chu et al.24 who reported a room temperature  Young’s modulus of 360 GPa and Poisson’s ratio of 0.189. Pastor and Meyer22 first studied the effect of addition of TaSi2 and other silicides on the densification and oxidation resistance of ZrB2,  assessing the formation of boride-silicide solid solutions and improvement of oxidation resistance in air up to 12001C. Recently, Talmy et al.25 have studied ceramics in the system ZrB2-Ta5Si3 and have reported a significant improvement in the densification, which was fully accomplished at 19001C and of the oxidation resistance in  comparison with pure zirconium diboride. Furthermore, these au thors have detected the formation of a ZrB2-based solid solution phase, due to tantalum entering the diboride lattice. Opila et al.26  studied the influence of TaSi2 additions to ZrB2-SiC compositions and found a significant improvement in the oxidation resistance at 16271C in air. Finally,  it was reported that TaSi2 is beneficial as a sintering additive (2 vol%) for a HfB2-based material,27 due to  liquid phase formation.  In this  contribution,  the densification, microstructure and  properties of ZrB2 and HfB2 composites containing 15 vol% TaSi2 are presented and discussed. The volumetric fraction of the TaSi2 addition was chosen to allow for direct comparison of these composites to previously produced UHTC composites  where the matrix phases were the same but the secondary phase was MoSi2.20,21  II.  Experimental Procedure  Commercial powders were used to prepare the ceramic compos ites:  hexagonal  ZrB2 impurities (maximum content): C: 0.25 wt%, O: 2  grade  B  (H.C.  Starck,  Karlsruhe,  Germany),  wt%, N: 0.25 wt%, Fe: 0.1 wt%, Hf: 0.2 wt%, particle size range 0.1-8.0 mm; hexagonal HfB2 waukee, WI), \\x00325 mesh, particle size range 0.5-5.0 mm, mean (Cerac Incorporated, Milparticle size B1 mm (by scanning electron microscopy (SEM)  analysis),  impurities: Al (0.07%), Fe (0.01%), Zr (0.47%), and  hexagonal TaSi2 (ABCR, GmbH & Co, Karlsruhe, Germany), \\x0045 mm. The following compositions were prepared: A. ZrB2115 vol%TaSi2, labeled ZBT. HfB2115 vol%TaSi2, The powder mixtures were ball milled for 24 h in absolute  B.  labeled HBT.  ethanol using zirconia milling media. Subsequently the mixtures  were dried in a rotary evaporator and sieved through a 60-mesh  screen. Some preliminary tests were carried out  to attempt  to  pressurelessly sinter these materials but similar to results reported in the literature22 compositions with 10 or 15 vol% TaSi2 contained significant amounts of porosity (B10%) as ascer tained by SEM analysis. Hot pressing was, therefore, chosen to  densify the TaSi2-contaning compositions. Hot pressing was conducted in low vacuum (B100 Pa) using an induction-heated  graphite die with a constant uniaxial pressure of 30 MPa, heating rate 201C/min and free cooling. For each composition, the  maximum sintering temperature was set on the basis of the i.e., 18501C/10 min and 19001C/15 min for  shrinkage  curve,  ZBT and HBT, respectively. The bulk densities were measured  by the Archimedes’ method. Crystalline phases were identified  by  X-ray  diffraction  (XRD)  (Siemens D500, Karlsruhe,  Germany). The microstructures were  analyzed using  (SEM,  Cambridge S360, Cambridge, U.K.) and energy dispersive spec troscopy  (EDS,  INCA Energy  300,  Oxford  instruments,  Oxford, U.K.) on polished surfaces. Mean grain sizes, amount  of porosity and of secondary phases were determined through  image analysis on SEM micrographs of polished surfaces using a  R. Cutler—contributing editor  w  Author to whom correspondence should be addressed. e-mail: diletta.sciti@istec.cnr.it  Manuscript No. 24398. Received March 11, 2008; approved May 28, 2008.  Journal  J. Am. Ceram. Soc., 91 [10] 3285 - 3291 (2008)  DOI: 10.1111/j.1551-2916.2008.02593.x  r 2008 The American Ceramic Society  3285  \\x0c', 'commercial  software program (Image Pro-plus  4.5.1, Media  Cybernetics, Silver Springs MD).  Vickers microhardness (HV1.0) was measured with a load of  9.81 N, using a Zwick 3212 tester  (ULM, Germany). Young’s  modulus (E) was measured by the resonance frequency method on 28 mm \\x02 8 mm \\x02 0.8 mm specimens using an HP gain-phase analyzer (Tokyo, Japan). Fracture toughness (KIc) was evaluated using chevron-notched beams (CNB) in flexure. The test bars, 25 mm \\x02 2 mm \\x02 2.5 mm (length by width by thickness, respectively), were notched with a 0.1 mm-thick diamond saw; the chev ron-notch tip depth and average side length were about 0.12 and  0.80 of  the bar thickness, respectively. The specimens were frac tured using a semiarticulated silicon carbide four-point fixture with  a lower span of 20 mm and an upper span of 10 mm using a screw driven load frame (Instron mod. 6025, High Wycombe, U.K.). The  specimens, three for each composite, were loaded with a crosshead  speed of 0.05 mm/min. The ‘‘slice model’’ equation of Munz et al.28 was used to calculate KIc. On the same machine and with the same fixture, the flexural strength (s) was measured at room temperature, 12001 and 15001C in air on chamfered bars 25 mm \\x02 2.5 mm \\x02 2 mm (length \\x02 width \\x02 thickness, respectively), using a crosshead speed of 0.5 mm/min. For the high-temper ature tests, a soaking time of 18 min was set  to reach thermal  equilibrium. Five specimens were tested at room temperature, three at 12001 and 15001C.  III.  Results  (1)  Densification  ZBT started shrinking at 15801C (Fig. 1) and required a maximum temperature of 18501C to achieve a nearly full density  with a dwell time of 10 min. Its final bulk density was 6.6 g/cm3  and a mass loss of about 15% was measured. For HBT, the mixture started shrinking at 17301C and the densification was completed at 19001C, with a holding time of 15 min (Fig. 1). The final bulk density was 10.7 g/cm3 with a mass loss of 8%.  During all of the sintering cycles, a significant variation of  the  pressure inside the furnace vacuum chamber was observed. At T412001C,  the pressure  increased from the  initial  value of  60 Pa up to a maximum of 130 Pa, which occurred in the temperature range 13001-14501C. In the temperature range 14501- 16501C, the pressure decreased reaching the initial value around 16801-17001C. An example of pressure variation is shown in the  inset  in Fig. 1. The reason for this pressure variation was  the  release of gaseous species, as explained in Section IV.  (2)  XRD Analysis  A. ZBT: Hexagonal ZrB2, hexagonal TaSi2 and traces of tetragonal ZrO2 were detected along with a series of reflections that were attributed to a solid solution formed by Ta dissolution into  the ZrB2 lattice (Fig. 2(a)). Compared with pure ZrB2, peaks were shifted towards higher angles, which indicates a  these  contraction of the diboride unit cell. The unit cell parameters of this newly formed phase were a 5 3.152 A˚ and c 5 3.485 A˚ , i.e., (a 5 3.169A˚ , c 5 3.530 A˚ ).  shorter  than those of pure ZrB2 shift was more pronounced at higher diffraction angles,  The  as  shown in the diffraction pattern of Fig. 2(b). On the basis  of the Vegard’s rule, for the system ZrB2-TaB2 and hypothesizing that only Ta can enter the ZrB2 structure, it can be estimated that the amount of Ta incorporated into ZBT was about  18-20 at.%, giving (Zr0.8Ta0.2)B2 as the composition of the solid solution.  B. HBT: The XRD pattern indicated (Fig. 3) that hexagonal  HfB2, hexagonal TaSi2 and monoclinic HfO2 were the crystalline phases. The diffraction spectrum is reported in Figs. 3(a)  and (b). At high diffraction angles, splitting of the main reflec tions of HfB2 was observed. Different from ZBT, the secondary peaks for HBT were visible only at high diffraction angles (2y4601). These additional reflections were identified as a solid  solution formed by the incorporation of Ta into the HfB2 lattice. The unit cell parameters of this newly formed phase were a 5 3.131 A˚ and c 5 3.440 A˚ , HfB2 (a 5 3.140 A˚ , c 5 3.470 A˚ ). According to the Vegard’s rule applied to HBT, the amount of Ta incorporated into HBT was  i.e.,  shorter  than those of pure  about 18-20 at.%, giving (Hf0.8Ta0.2)B2 as the composition of the solid solution.  For both the composites, the nonuniform Ta distribution was  due to the fact that the processing time was not sufficient to al low for homogenization under the conditions used for densifi cation. This feature is in agreement with the findings of Talmy et al.25 for ZrB2-Ta5Si3 composites.  -0.5  0.5  1.5  2.5  3.5  4.5  5.5  0  t (min)  S  h  r  i  n  k  e g a  (  m m  )  1000  1100  1200  1300  1400  1500  1600  1700  1800  1900  2000  2100  T  ˚ (  C  )  ZrB2  HfB2  Ton:1580˚C  Ton:1730˚C  30  25  20  15  10  5  Fig. 1.  Shrinkage behavior and temperature profile of the composites.  Inset: pressure variation inside the hot-pressing chamber as a function of  the temperature.  90  ZrB2 Zr-B-Ta ss  2-Theta°  20  ZrB2 TaSi 2 Zr-B-Ta ss m-ZrO2  I  n  t  n e  s  i  t  y  (  a  .  u  )  2-Theta°  70  60  50  40  30  150  140  130  120  110  100  (a)  (b)  Fig. 2.  X-ray diffraction spectra of  the ZrB2-15 vol% TaSi2 composite.  (a) 2y: 201-701,  (b) 2y: 881-1581. Reflections  from CuKa2  radiation were  removed.  3286  Journal of the American Ceramic Society—Sciti et al.  Vol. 91, No. 10        \\x0c', 'October 2008  Sintering and Mechanical Properties of ZrB2-TaSi2 and HfB2-TaSi2 Ceramic Composites  3287  (a)  )  .  u  .  a  (  y  t  i  s  n e  t  n  I  HfB2 TaSi 2  Hf-B-Ta ss HfO2  (b)  HfB  2  Hf-B-Ta ss  20  30  40  50  60  70  94  96  98  100  102  104  106  2-Theta˚  2-Theta˚  Fig. 3.  X-ray diffraction spectra of  the HfB2-15 vol% TaSi2 composite.  (a) 2y: 201-701,  (b) 2y: 941-1061. Reflections from CuKa2  radiation were  removed.  (3) Microstructural Features  A. ZBT: Some examples of polished surfaces of ZBT are shown in Figs. 4(a)-(c). Little or no porosity (o1%) was observed, dicating that the relative density was higher than 99%. The  in globular ZrB2 grains visible in Fig. 4(a) had an average size of about 2 mm. These grains were surrounded by a Zr-Ta-B solid  solution, which appeared lighter in color (Fig. 4(b)). According  to quantitative EDS analysis, the composition of  this solid so lution was  (Zr0.8Ta0.2)B2 predicted by the Vegard’s law. Pure TaSi2 was not clearly identified in the composite, but a TaxSiy phase containing a certain amount of Zr was observed (Fig. 4(d)). This phase had a lighter  in agreement with the  composition  color than ZrB2 and (Zr,Ta)B2. By image analysis, the content of the TaxSiy phase was estimated to be 3-4 vol%. Residual ZrO2 particles were also found along with silica-based phases containing various impurities. These oxide phases were attrib uted to the oxygen impurities  in the starting powders and/or  oxygen contamination during powder processing. Spurious car bide phases, such as Zr—Ta-C, SiC, or Si-C-O, were also de tected in limited amounts  (Fig.4(c)). Carbon impurities were  certainly present in the starting ZrB2 powder but some carbon enrichment could have also resulted from the polyethylene bot tles used for the milling procedure. Moreover,  in the graphite rich sintering environment, CO generation could have induced  either the carburization of metals or the carbothermal reduction  of oxide species, as observed for other composites of class of materials.29,30 B. HBT: A fine microstructure with little porosity (o1%) was observed in HBT in Fig. 5(a). The mean HfB2 grain size was around 1 mm, which was similar to the starting particle size of  the same  the HfB2 powder, indicating that no significant grain coarsening occurred during sintering. The lightest colored phase was iden in agreement with the findings of  the X-ray  tified as HfO2 diffraction pattern. Significant  fractions of HfO2 were in other composites produced from the same starting HfB2 powder as the present work,19,20 suggesting a large oxygen contam found  ination. TaSi2 had a very irregular shape (Fig. 5(b)) and the way this phase filled the spaces between the HfB2 grains indicated possible high-temperature ductile behavior. Other  spurious  phases were mixed carbides, based on (Hf, Ta)C and SiC, whose  Fig. 4.  Polished surfaces of  the ZBT.  (a) Overall view,  (b) enlarged view showing pure diboride grains  (1)  surrounded by the solid solution (2).  (c) Carbide grain. (d) EDS spectra related to the phases labeled in (b) and (c). Legend: (1) ZrB2, (2) (Zr,Ta)B2 (3) TaSi2, (4) (Zr,Ta)C. Beam accelerating voltage: 15 keV.    \\x0c', '3288  Journal of the American Ceramic Society—Sciti et al.  Vol. 91, No. 10  Fig. 5.  Polished surfaces of the HBT. (a) Overall view, (b) and (c) details of the microstructure, (d) EDS spectra related to the phases labeled in (b) and  (c). Legend:  (1) TaSi2,  (2) Ta-Hf-C,  (3) SiC-based phase. Spectra 1 and 2 collected at 15 keV, spectrum 3 collected at 6 keV to limit beam lateral  spreading.  morphologies were  consistent with  a  liquid-phase  sintering  formation of a solid solution was more pronounced in ZBT than  mechanism. The origin of  these  carbides may be ascribed to  in HBT.  the interaction of  the sintering environment with the starting  powders, as previously mentioned for ZBT. Analyzing the mi crostructure  in BSE imaging, many HfB2 grains core-shell structure (Fig. 6). By EDS analysis, the outer shell was  exhibited a  estimated to be a solid solution with composition (Hf0.8 Ta0.2)B2 in agreement with the composition predicted by the Vegard’s  law. The microstructural and XRD analyses  showed that  the  (4) Mechanical Properties  The mechanical properties are reported in Table I, along with  some data  for  two other  composites with the  same matrix  phases, i.e., HfB2 and ZrB2, but with MoSi2 as the secondary phase.20,21 Generally speaking, HBT had higher hardness, stiff ness, and fracture toughness, but lower room-temperature flexural strength than ZBT. Both at 12001 and 15001C, the strength  of HBT (700 and 600 MPa, respectively) was higher than ZBT  (600 and 370 MPa,  containing composites,  respectively). With respect  to the MoSi2it can be seen that TaSi2 additions increased hardness and toughness but had almost no effects on Young’s modulus. The flexural strength of ZBT (840733 MPa)  was significantly higher same matrix with MoSi2 (704798 MPa). On the other hand, the strength of HBT (698758) and of the HfB2-MoSi2 (7427142) were not differfrom a statistical point of view, but in the latter case the  than  the  ent  values were affected by a very large data dispersion.  IV.  Discussion  During the densification of ZBT and HBT, several phenomena  occurred,  the most  important being the mass  loss and partial  substitution of Ta for Zr or Hf,  implying that Ta has at  least  limited diffusivity in both borides. The mass loss can be attrib uted to the interaction of TaSi2 either with B2O3, which is present as surface oxide on the boride particles, or with CO  generated inside the graphite-rich environment of  the furnace.  The pressure in the hot press vacuum chamber increased above 13001C due to the formation of volatile species and reached its maximum (130 Pa) around 14001C. Using a commercial  soft ware package  (HSC Chemistry for Windows 5, Outokumpu  Research Oy, Pori, Finland), some potential reactions between  TaSi2 and CO, TaSi2, and B2O3 were analyzed in the range 12001-19001C under a pressure of 100 Pa (i.e., approximately  Fig. 6.  Polished surface of HBT. Enlarged view showing pure diboride  grains (1) surrounded by the solid (Hf,Ta)B2 solution (2).  \\x0c', 'the vacuum level inside the hot-press chamber), considering pure  materials (i.e., no solid solutions), with the activity of each phase  equal to one. For simplicity, the molar ratio between the phases  was  taken equal  to one. The thermodynamic results are sum marized in Figs. 7(a) and (b) and give the following indications:  (1) CO promotes TaSi2 decomposition with the formation of T \\x15 13001C T \\x15 14501C. SiO T \\x15 12001C, the formation of TaC and SiC is expected. In the gas at and Si(g) at For range 14501-18501C, the formation of liquid Si is predicted. The from \\x000.9 \\x01 103 kJ at  change in the Gibbs free energy varies 12001C to \\x001.3 \\x01 103 kJ at 19001C. (2) B2O3 promotes TaSi2 decomposition with the formation at T \\x15 12001C and Ta of SiO gas at T412001C, TaB2 at T \\x15 17001C. At T \\x15 14001C, the emission of BO also occurs. For this reaction the change in the Gibbs free energy varies from \\x001.9 \\x01 103 kJ at 12001C to \\x002.6 \\x01 103 kJ at 19001C. The above thermodynamic predictions are in partial agree ment with the SEM observations. TaSi2 decomposition and consequent SiO volatilization was certainly confirmed by  weight  losses,  the increase in pressure in the furnace cham ber, XRD data and microstructural analyses. No Ta or TaB2 species were detected either by XRD or by SEM analy sis. However,  these  species  are  expected to form the  solid  solution with the borides. There was also evidence of carbide  formation, which were mixed phases of Ta and Zr (or Hf). In  both ZBT and HBT, Si-based phases with a very irregular  morphology suggesting liquid-phase behavior were often de tected. Possible origin of  these phases  can be  either  carbo thermal reduction of residual silica or carburization of liquid  silicon.  As far as densification is concerned,  therefore,  the two key  mechanisms are the formation of solid solutions and Sior SiO2based liquid phases.27 Due to the significant TaSi2 decomposition observed in ZBT, it is believed that in this composite  the  densification was  primarily  governed  by  the  solid  so lution formation.  In the case of HfB2 both mechanisms were  equally active.  The presence of  solid solutions with unknown properties  complicates the analysis of  the mechanical properties. The fol lowing discussion will therefore focus on the properties of well characterized phases. HBT was harder  than ZBT because the  HfB2 matrix is harder than the ZrB2 matrix. To give reference values, the Knoop hardness measured with a load of 160 g is about 23.5 GPa for HfB2 and 20.5 GPa for ZrB2.31 The hardness of the composites was lower than the pure matrices as the hardness of TaSi2 is about 15.6 GPa.23 Because the hardness of MoSi2 is about 12 GPa,32 the TaSi2-reinforced composites were harder than the MoSi2-reinforced composites. The higher Young’s modulus of HfB2 than ZrB2 made HBT stiffer than ZBT.33 Also in this case, the presence of TaSi2 lowered the the composites relative to pure matrices.  some  Young’ modulus of  The elastic modulus of the composites was not sensitive to the  change of secondary phase, although MoSi2 has a higher stiffness (E 5 425 GPa34) than TaSi2. One possible explanation for this behavior is that the MoSi2-reinforced composites contained a not negligible content of silica (4%-5%) which decreased the stiffness.21 HBT was about 10% tougher than ZBT, whilst the  room-temperature strength of ZBT was about 20% higher than  HBT. The reasons for this crossover are beyond the scope of this  study and may be  the  focus of a future work. The flexural  strength was  the property most  influenced by the type of  the  secondary phase. Replacing MoSi2 with TaSi2, in fact, increased the strength of the ZrB2-based composite and reduced the data dispersion of the HfB2-based composite. Fractographic analysis of specimens fractured at room temperature showed that in the  MoSi2-based composites, large MoSi2 agglomerates were present on fracture surfaces.18,20 No TaSi2 agglomerates were instead detected on the fractured surfaces of ZBT and HBT,  implying that TaSi2 dispersed more uniformly in the matrices than MoSi2. At elevated temperatures, the strength of HBT was higher than ZBT. A similar crossover was observed for the same  matrices containing MoSi2, indicating that the overall strength may be dictated by the matrix. At 12001C the load-displacement curves of ZBT and HBT were linear up to fracture. At 15001C a  Table I.  The Mechanical Properties of the Tested Composites in Comparison with the Same Matrices Reinforced with MoSi2  Composition  Label  Density  HV  E  KIc  sRT  s12001C  s15001C  Experimental  Relative  Vol%  g/cm3  %  GPa  GPa  MPa \\x01 m1/2  MPa  MPa  MPa  ZrB2-15 TaSi2 HfB2-15 TaSi2 ZrB2-15 MoSi2 HfB2-15 MoSi2  ZBT  6.6  99  17.870.5 21.970.5 14.970.5 20.670.4  444724 52875 45274 53075  3.870.1 4.170.1 3.570.6 3.870.1  840733 698758 704798 7427151  598725 703724  37475 597746 333731 548720  HBT  10.7  99  w  —  6.0  99  — 664728  z  —  10.2  99  Mean71 standard deviation.  w  Sciti et al.21;  z  Sciti et al.20  -0.2 1100  0.0  0.2  0.4  0.6  0.8  1.0  1.2  1.4  1.6  1.8  2.0  1 mol TaSi 2+ 1 mol B2O3  Ta2O5  BO  B2 O3  SiO2  TaB 2  Ta  SiO (g)  R  a e  c  t  i  n o  p  r  u d o  c  t  s  (  m  o  l  )  Temperature (˚C)  0.0  0.2  0.4  0.6  0.8  1.0  (a)  (b)  CO  1 mol TaSi 2+ 1 mol CO(g)  residual TaSi 2  Si (g)  SiO(g)  TaC  SiC  Si (l)  R  a e  c  t  i  n o  p  r  u d o  c  t  s  (  m  o  l  )  Temperature (˚C)  2000  1900  1800  1700  1600  1500  1400  1300  1200  1100  2000  1900  1800  1700  1600  1500  1400  1300  1200  Fig. 7. Molar content of the products deriving from the reactions: (a) 1 mol TaSi211 mol CO(g), (b) 1 mol TaSi211 mol B2O3, as a function of the temperature and at constant pressure of 100 Pa.  October 2008  Sintering and Mechanical Properties of ZrB2-TaSi2 and HfB2-TaSi2 Ceramic Composites  3289          \\x0c', '3290  Journal of the American Ceramic Society—Sciti et al.  Vol. 91, No. 10  Fig. 8. (a), (b) Examples of fracture surface morphology (detail) and cross section after the 15001C test of a ZrB2-15 vol%TaSi2 specimen, (c) Fracture surface after the 12001C test of a HfB2-15 vol% TaSi2 specimen, (d) HfB2-15 vol% TaSi2 specimen surface after the 15001C test showing a dense glassy layer containing HfO2 crystals.  slight departure from linearity was observed only for HBT. This  temperature the flexural strength of HBT was higher than ZBT.  suggests that softening of the secondary phase is not likely to be  the main reason for  the high-temperature  strength decrease.  More  likely,  it was  the  combination of  the high-temperature  fracture toughness and the oxidation attack that determined the  Comparing the same matrices with MoSi2 in place of TaSi2, was shown that TaSi2 increases the room-temperature properties and the high-temperature strength. Remarkably, the rethe HBT composite at 15001C (600 MPa)  tained strength of  it  high-temperature strength behavior. Both composites,  in fact,  was about 86% of the room-temperature value.  underwent significant surface modification at  the test tempera tures (Fig. 8), due to oxidation. Glass formation can either blunt  cracks  leading to strength retention or  introduce new defects  causing strength degradation. Work on quantifying oxidation resistance is in progress. The retained strength of ZBT at 15001C  was about 45% of  the room-temperature value, but  for HBT  this ratio was 86%. Among borides, strength retention at ele vated temperature has only been approached by a HfB2-20 vol% SiC composite densified by spark plasma sintering.5,27  V.  Conclusions  ZrB2 and HfB2 plus 15 vol% TaSi2 additions were densified to 99% of relative density by hot pressing in the temperature range 18501-19001C. For both compositions, gaseous phases were re leased during sintering and mass  losses were measured after  sintering. The final microstructures of both composites were al(1-3 mm). The most  most pore-free with small grain sizes  im portant microstructural features were the partial decomposition  of TaSi2 and the formation of solid solutions in the (Zr,Ta)B2 and (Hf,Ta)B2 systems. By thermodynamic analysis, the decomposition of TaSi2 was attributed either to the interaction with CO species which form in the graphite-rich reducing sintering  environment, or to the interaction with B2O3 present as oxide impurity on boride particle surfaces.  The HfB2-based composite was harder (22 GPa), stiffer (528 GPa), and tougher (4.1 MPa m1/2) than the ZrB2-based composite. However, the latter had a higher room-temperature strength (840733). On the other hand, when tested at elevated  Acknowledgments  The authors wish to thank Daniele Dalle Fabbriche for the hot pressing routes.  One of  the authors (L. S.) gratefully acknowledges the financial support of  the  Regional Project MATMEC.  The authors would like to thank the referees and the associated editor for their  useful and constructive suggestions for this work.  References  1C. Mroz, ‘‘Zirconium Diboride,’’ Am. Ceram. Soc. Bull., 73, 141-2 (1994). 2R. A. Cutler, ‘‘Engineering Properties of Borides’’; pp. 787-803 in Ceramics and  Glasses: Engineered Materials Handbook, Vol. 4, Edited by S. J. Schneider Jr.  ASM International, Materials Park, OH, 1991. 3M. Pastor,  ‘‘Metallic Borides: Preparation of Solid Bodies. 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Lorincz, ‘‘Oxidation of ZrB2 and HfB2-Based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mat. Sci., 39, 5969-  77 (2004). 27F. Monteverde,  ‘‘Ultra-High Temperature HfB2-SiC Ceramics Consolidated by Hot Pressing and Spark Plasma Sintering,’’ J. Alloys Compounds, 428, 197-205  (2007). 28D. G. Munz, J. L. Shannon Jr., and R. T. Bubsey,  ‘‘Fracture Toughness Cal culation from Maximum Load in Four Point Bend Tests of Chevron Notch Spec imens,’’ Int. J. Fract., 16, R137 (1980). 29D. Sciti, L. Silvestroni, and M. Nygren,  ‘‘Spark Plasma Sintering of Zrand  Hf-Borides with Decreasing Amounts of MoSi2 as Sintering Aid,’’ J. Eur. Ceram. Soc., 28, 1287-96 (2008). 30R. M. German,  in Sintering Theory and Practice, p. 432. John Wiley & Sons,  Inc, New York, 1996. 31J. F. Shackelford and W. Alexander, CRC Materials Science and Engineering  Handbook. CRC Press, Boca Raton, Fl.; London, 2001. 32D. Sciti, S. Guicciardi, and A. Bellosi,  ‘‘Microstructure and Properties of  Si3N4-MoSi2 Composites,’’ J. Ceram. Proc. Res., 3 [3] 87-95 (2002). 33N. P. Bansal, Handbook of Ceramic Composites. Kluwer Academic Publishers,  Boston, 2005. 34M. Nakamura, S. Matsumoto, and T. Hirano,  ‘‘Elastic Constants of MoSi2 and WSi2 Single-Crystals,’’ J. Mater. Sci., 25 [7] 3309-13 (1990).  &  \\x0c']"
},{
  "_id": 243,
  "PDF": "Sintering and oxidation behavior of HfB2-SiC composites from 0 to 30 vol_ SiC between 1450 and 1800 K.pdf",
  "Text": "[\"Ceramics International 45 (2019) 1846-1856  Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Sintering and oxidation behavior of HfB2-SiC composites from 0 to 30 vol% SiC between 1450 and 1800 K  T  Cassandre Pirioua, Olivier Rapauda,⁎, Sylvie Foucauda, Ludovic Charpentierb, Marianne Balat-Pichelinb, Maggy Colasa  a IRCER-CNRS, UMR 7315, Centre Européen de la Céramique (CEC), 12 Rue Atlantis, F-87068 Limoges Cedex, France b PROMES-CNRS, UPR 8521, 7 Rue du Four Solaire, F-66120 Font-Romeu Odeillo, France  A R T I C L E  I N F O  A B S T R A C T  Keywords: Composites Sintering Corrosion  Fully-dense HfB2-SiC composites with controlled microstructure were obtained by Spark Plasma Sintering for 0-30 vol% SiC. Samples were then oxidized in a solar furnace at 1450, 1550 and 1800 K for 20 min under stagnant air. To complete this approach, the oxidation kinetics were followed by thermogravimetric analysis (TGA) for isothermal runs using the same conditions. Oxidized samples were characterized by X-ray diffraction, scanning electron microscopy and Raman mapping. The highest oxidation resistance, characterized by no mass evolution after 7 min exposure at 1800 K, was obtained for the sample containing 20 vol% SiC. At the same temperature, the oxidation kinetics of a composite with a lower amount of SiC (typically 10 vol%) was controlled by the diffusion of oxygen up to 10 µm in depth whereas a sample with 30 vol% SiC showed mass losses after 5 min exposure proving the presence of a non-protective oxide layer due to vaporization. Blisters of a glassy phase surrounded by hafnia were observed at the surface of all oxidized samples at 1800 K.  1.  Introduction  In the aeronautical industry, vehicles have to endure oxidizing atmospheres at temperatures up to 1800 K. This concerns in particular the structural materials constituting the combustion chamber of the engine [1]. Therefore, materials have to possess a high melting point and a good oxidation resistance up to 1800 K. Borides and carbides of the IVb group are promising materials because of their high melting points. However, due to the lower thermal conductivities of such carbides (< 40 W m−1 K−1 from room temperature to 1073 K [2]), diborides and more specifically zirconium and hafnium diborides (ZrB2 and HfB2) are the most attractive candidates of the IVb group. Nevertheless, these diborides cannot be used alone because of the volatilization of B2O3(g) during oxidation which leaves behind a porous layer of their oxides above 1373 K (HfO2(s), ZrO2(s)). Previous works have shown the efficiency of adding SiC to HfB2 and ZrB2 in order to improve the oxidation resistance by the formation of a protective glassy layer at the surface [2-4]. It has been shown that HfB2 and its oxide HfO2 possess higher melting points and better oxidation resistance than ZrB2 and ZrO2 [5,6]. Moreover, Young's Modulus and the thermal conductivity are higher for HfB2 (530 GPa [7] and 105 W m−1 K−1 at 293 K [8]) than ZrB2 (489 GPa [5] and 84 W m−1 K−1 at 293 K [8]). All these properties  make HfB2 the most interesting material for such applications. Thermodynamically, HfB2 and SiC should react with oxygen and form oxidized species depending on the temperature and oxygen partial pressure. Oxidation reactions of HfB2 and SiC are the following:  HfB  2 ( s  )  +  SiC  ( s  SiC  ( s  )  )  +  +  3 2  O  5 2  O  2 ( g  )  =  HfO  2 ( s  )  +  B O  2  3 ( g  O  2 ( g  )  =  SiO  2 ( ) l  +  CO  ( g  )  2 ( g  )  =  SiO  ( g  )  +  CO  ( g  )  )  (1)  (2)  (3)  Balat et al. [9] have shown for bulk SiC samples that the oxidation depends on the temperature, the oxygen partial pressure but also on the SiC crystalline structure and the oxidizing atmosphere. Passive and active oxidation reactions are presented in Eqs. (2) and (3), respectively. Several authors have pointed out the active oxidation of SiC in depth of composites with HfB2 or ZrB2 when the oxygen partial pressure is very low (PO2  10−1 to 10−15 Pa) [10,11]. Therefore the volatilization of SiO(g) should leave behind a porous structure, which is often called a SiC depletion layer. Moreover, whereas some authors [2,7] explained that the as-formed products (B2O3(l) and SiO2(l)) generate the formation of a protective borosilicate glass in a higher temperature  ⁎ Correspondence to: IRCER, Centre Européen de la Céramique, 12 rue Atlantis, F-87068 Limoges Cedex, France. E-mail address: olivier.rapaud@unilim.fr (O. Rapaud).  https://doi.org/10.1016/j.ceramint.2018.10.075 Received 1 June 2018; Received in revised form 8 October 2018; Accepted 8 October 2018  Available online 12 October 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  \\x0c\", 'C. Piriou et al.  Table 1  Summary of compositions, processing and relative densities.  Ceramics International 45 (2019) 1846-1856  HfB2 content (vol%) SiC content (vol%) Imposed temperature (K) Relative density (%)  H100  100 0 2123 100  H95S5  95 5 2123 99  H90S10  H85S15  H80S20  H75S25  H70S30  90 10 2073 100  85 15 2023 100  80 20 2023 100  75 25 1973 99  70 30 1923 99  range than SiO2(l) or B2O3(l) glasses alone, others [4] highlighted the total volatilization of B2O3(g) and thus, the presence of a silica glass only at the surface. Several authors [12-14] also disagreed about the optimal composition between 15 and 30 vol% SiC added to HfB2 in terms of oxidation resistance. It has to be noted that the oxidation behavior can depend on the experimental method of sintering (Hot Pressing, Spark Plasma Sintering) and its parameters (temperature, holding time, pressure, etc). Some other parameters such as powder granulometry and distribution of phases can also impact the microstructures and the relative densities of the final materials. It is obvious that a low relative density decreases the oxidation resistance because of the rapid diffusion of oxygen through the pores. It has been shown that particle sizes of 5 µm and 1 µm for HfB2 and SiC, respectively, encouraged densification [5]. The discrepancies of the optimal composition can also be related to the technique of oxidation (electric arc furnace, oxyacetylene torch, air plasma, solar furnace, thermogravimetric analysis under air) and its parameters (temperature, holding time, pressure, presence of impurities in the atmosphere, etc). Therefore, the wide range of experimental conditions can lead to different results in terms of oxidation behavior. So far, no study has been published showing the most resistant to oxidation composition from 0 to 30 vol% SiC in the HfB2-SiC system by controlling accurately both microstructure (grain size, secondary phases and phase distribution), and relative density (> 99%) of each sample. Moreover, oxidation kinetics of all these compositions have always been studied with thermogravimetric analysis using slow heating rates from room temperature. This will have an influence on the oxidation mechanisms. Such equipment is not able to elevate the temperature as fast as in the considered aeronautical applications where the transient heating regimes are shorter than one minute. In this work, the main objective is to study precisely the effect of the silicon carbide added to the hafnium diboride on the oxidation resistance by controlling microstructures and relative densities of the starting materials. To carry out this goal, several compositions from 0 vol% to 30 vol% SiC were sintered. Fine powder granulometries centered around 1 µm and 5 µm in diameter for the silicon carbide and hafnium diboride were respectively chosen. Moreover, an optimization of the sintering parameters was carried out for every composition in order to obtain fully-dense materials with a controlled grain size. Oxidation in a reactor placed in a solar furnace was used in order to reach temperatures above 1450 K in a very brief time (less than 1 min) and therefore the materials can be considered as subjected immediately to high temperature. Samples were oxidized in the REHPTS facility available at PROMES, France, at three different temperatures: 1450, 1550 and 1800 K. These temperatures have been selected referring to the work by Monteverde and Bellosi [3] who described the formation of a protective borosilicate glass at the surface of the material under oxidation. Therefore, these three temperatures are characteristic of the main steps of the oxidation mechanisms. The final goal of this paper is thus to achieve a better understanding of the influence of the composition on the oxidation behavior. The crystalline phases before and after oxidation were investigated by X-ray diffraction. The expected glassy phase at the surface was characterized by Raman mapping in a completely new way in order to get both chemical information and spatial location. A complementary and innovative work was carried out with a thermogravimetric analyzer with hot insertion of samples, which means  bypassing any contribution of ture.  the heating ramp from room tempera 2. Experimental procedure  2.1. Sample preparation  ρ = 11 g cm−3, abcr, HfB2 powder (99.5% purity, d50 = 22 µm, (α-SiC, > 98.5% purity, Germany) and SiC powder d50 = 1 µm, ρ = 3.2 g cm−3, H.C. Starck, Germany) were chosen as raw materials. For HfB2, Zhu et al. [15] described the utility of decreasing the diboride grain size in order to enhance sintering with SiC. To this end, the HfB2 powder was planetary milled using 5 cycles of 6 min split into 1 min milling followed by 5 min break to prevent heating. A WC media with WC balls was used and a median diameter of 5 µm was achieved for the HfB2 grains. Different amounts of SiC were added to the HfB2 powder and mixed to obtain a homogeneous distribution of both phases (Table 1). Mixtures were weighed to obtain 3 mm thickness pellets after sintering, pre-compacted at 75 MPa in a graphite die (inner diameter 13 mm), lined with a graphitized paper (Papyex, Mersen, France) and sintered using Spark Plasma Sintering (Dr Sinter 825, Fuji Electronics Industrial Co. Ltd., Japan). The temperature was monitored by an optical pyrometer at the external surface of the die. Powders were sintered at different temperatures, according to the SiC content (Table 1). Every sintering cycle was carried out in vacuum, using a heating rate of 100 K min−1 with an applied load of 100 MPa and 10 min of holding time. A DC current with a pulse sequence of 12 ms on (ton) and 2 ms off (toff) was used for heating. The imposed temperatures during sintering for every composition were selected after an optimization which will be subsequently discussed. The sintered bodies were labelled using a code based on the HfB2 and SiC contents. The H and S letters, attributed to the HfB2 and SiC phases respectively, were followed by the volume percent of HfB2 and SiC constituting the composite. Quantitative carbon and oxygen analyses were performed on HfB2 powder before and after milling using the Emia-321V (Horiba, Japan) carbon analyzer and the EMGA-830 (Horiba, Japan) oxygen analyzer. X-ray fluorescence (WD-XRF ARL OPTIM’X, Thermo Fischer, Madison, WI, USA) was used to characterize powders before and after milling in order to control the presence of impurities of the commercial and milled powders. The bulk density was measured using the method based on Archimedes’ principle with ethanol as the immersing medium. The relative density of every sintered material was evaluated by the ratio between the measured density and the theoretical density of powders calculated from the initial compositions using the rule of mixtures. Microstructures of sintered samples were analyzed using scanning electron microscopy (SEM; LEO 1530VP, Zeiss, Germany) combined with energy dispersive spectroscopy (EDS; Oxford, United Kingdom) for qualitative chemical analysis. The HfB2 grain size was measured on polished surfaces from SEM images using the ImageJ image processing software on at least 250 grains. Obtained values were corrected by a factor taking into account the tetrakaidecahedral shape of grains after sintering and the difference between the true and apparent sizes [16].  2.2. Oxidation tests  For the oxidation in the solar furnace, sintered bodies were polished  1847  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 1. X-ray diffraction patterns of: HfB2 commercial powder a), HfB2 milled powder b), as-sintered H80S20 c), oxidized H80S20 at 1450 K d) and oxidized H80S20 at 1800 K e).  with a diamond coated disk to remove the graphitized paper. They were then ultrasonically cleaned in acetone and ethanol and dried in air. All the compositions from 0 vol% to 30 vol% SiC (Table 1) of HfB2-SiC specimens in steps of 5 vol% SiC were exposed to air at 1450, 1550 and 1800 K for 20 min using the REHPTS reactor placed at the focus of the 6 kW solar furnace [17]. The samples were placed inside the reactor on a zirconia foam for thermal insulation from the water-cooled sampleholder. The concentrated solar flux was controlled by a shutter which adjusts the fraction of the flux radiating the sample. A monochromatic (5 µm) optical pyrometer measured the surface temperature. A relation for the temperature, which depends on the material emissivity and the optical path, was used to calculate the oxidation temperature after calibration with a blackbody. Samples were weighed before and after a given exposure time as well as the area of the exposed surface in order to obtain coherent mass gain rates expressed in mg cm−2 h−1. Four samples with similar composition were treated in the same conditions in order to evaluate the reproducibility of results. Therefore, standard deviation values were calculated and added to the mass gain rates. The oxidation kinetics were monitored by thermogravimetric analysis (TGA) using a unique isothermal SETARAM SETSYS EVO apparatus. Samples were placed in a home-made alumina crucible and heated isothermally in air at 1450, 1550 and 1800 K during 20 min. For oxidation by TGA, 4.5 × 4.5 × 3 mm3 samples were cut using a diamond saw and polished on all sides to obtain a mirror finish. Prior to oxidation by TGA, all the tested specimens were ultrasonically cleaned in acetone and ethanol, and then dried in air. The same compositions as those oxidized in the solar furnace were analyzed by TGA at the same temperatures for the same holding time. Moreover, three samples for each composition were oxidized with this apparatus and similar behaviors were obtained. With a view to simplification, not all the results will be presented. Crystalline phases were identified by X-ray diffraction (XRD; D8 Advance, Bruker, Germany) on starting powders, sintered and oxidized materials. Surfaces of sintered samples and cross-sections of the oxidized samples were polished up to a 1 µm finish. Microstructures of oxidized samples were analyzed using scanning electron microscopy (SEM; IT 300, Jeol, Japan) combined with energy dispersive spectroscopy (EDS; Oxford, United Kingdom) for chemical analysis. Raman mapping was performed to characterize the oxidized surfaces (InviaReflex, Renishaw, United Kingdom) by using the 532 nm wavelength at 1 mW on the sample for high resolution. To this end, a confocal mode and an objective ×100 were used and a spatial  resolution around 1 µm was achieved. The map was obtained by recording 6561 spectra on a surface of 40 × 40 µm2, with a step size of 500 nm. Nowadays, Raman spectrophotometers devoted to mapping provide a huge quantity of data. The Raman map extracted from the data can be built by classical methods. The principal one is called DCLS (Direct Classical Least Squares) and was achieved by extracting reference spectra from all the obtained data. Due to this approach, the vibration bands of Hf-O and Si-C, corresponding to hafnia and silicon carbide, respectively, were identified. By comparing the reference spectra of HfO2 and SiC with the experimental ones, new vibration bands were observed and associated to the Si-O bonds. However, classical methods are not appropriate in the case of such complex materials because of the overlapping of most of the bands. Thus, an accurate and more relevant method for complex materials, probably constituted of crystalline and vitreous phases, was used and is called PCA (Principal Component Analysis). This multivariate method is a decomposition procedure where a number of unknown orthogonal spectral sources (Principal Component, PC) are derived from data [18-21].  3. Results  3.1. Sintered samples  Little or no porosity was observed in these materials (Table 1). All samples achieved high densities above 99% of the theoretical density. XRD analyses on HfB2 powders before and after milling are presented in Fig. 1a) and b) and exhibit the presence of hexagonal HfB2, monoclinic HfO2 and cubic HfC in both cases. These last two secondary phases, also found in sintered HfB2-SiC bodies (H80S20 shown in Fig. 1c)), are likely to result as an excess product from the synthesis reaction used to produce hafnium diboride [22], corresponding to carbo/borothermal reduction of hafnia (Eq. (4)). This process often implies the presence of some HfC particles coming from the reaction between B4C and HfO2.  HfO  2 ( s  )  +  1 2  B C  4  ( s  )  +  3 2  C  ( s  )  HfB  =  2 ( s  )  +  2 CO  ( g  )  (4)  Quantitative carbon analysis revealed the presence of 0.04 wt% C in both commercial and milled HfB2 powders. Moreover, X-ray fluorescence (Fig. 2) detected the presence of Ta, Cu and W before and after milling which could be related to the impurities of starting powders  1848  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 2. Fluorescence spectrum of HfB2 powder before and after milling.  (HfO2, B4C and C) used for the synthesis of hafnium diboride powder (Eq. (4)). Therefore, the same amount of carbon before and after milling the W Lα1 fluoresgiven by quantitative analyses and the presence of cence line in both powders show that there was no tungsten carbide contamination during milling. The oxygen amount was determined by quantitative oxygen analysis and disclosed the presence of 0.299 wt% O and 0.405 wt% O in the commercial HfB2 powder and in the milled HfB2 powder, respectively. This provides evidence for a slight oxidation of powders, probably related to a slow increase of temperature during the milling. In the aim of simplification, the oxidation mechanisms will only take into account HfB2 and SiC. The SEM images of the obtained microstructures of sintered H90S10, H80S20 and H70S30 are presented in Fig. 3a)-c). Fig. 3d) is a higher magnification of sample H70S30 after sintering showing the primary and secondary phases. EDS spectra 1 and 2 reveal the presence of the primary phases HfB2 and SiC, respectively. Black grains assigned to SiC and grey grains to HfB2 are homogeneously distributed as it can be observed in Fig. 3a)-c). Secondary phases were also located: HfC, characterized by spectrum 3, was found at grain boundaries of the diboride phase and HfO2 inside the diboride grains (spectrum 4). Pores are indicated by arrows and as it can be seen in Table 1, confirm that little or no porosity is found in these materials. Fig. 4a) displays the evolution of the optimized sintering temperature as a function of the silicon carbide content. The sintering temperature can be clearly decreased when the amount of SiC increases. As the use of a lower temperature reduces grain growth, these results were used to optimize sintering parameters. Consequently, the imposed temperatures during Spark Plasma Sintering are presented in Table 1 for every composition. Fig. 4b) shows diboride grain sizes results after sintering of all compositions along with the values for the HfB2 milled powder for comparison purposes. The granulometry of the diboride powder (D50 = 5.4 µm) is close to the results obtained for HfB2 grains in H70S30 and H75S25. A little grain growth is observed for the H100 sample, but the median diameter only reaches 9.5 µm. Based on these similar microstructures (grain size after sintering and phase distribution) and the low values of porosity, oxidation behavior of all samples can be reliably compared.  3.2. Oxidized samples  The mass gain rate per unit time and per unit surface area for the samples oxidized in the REHPTS facility is plotted as a function of the temperature in Fig. 5. For the sample composed of 100 vol% HfB2 the  mass gain rate strongly increases with temperature, reaching 20 mg cm−2 h−1 at 1800 K. With the addition of SiC, the values attributed to the composites are close to each other and much lower with a maximum value of 3.2 mg cm−2 h−1 for the H95S5 sample oxidized at 1800 K. A little increase of mass gain rates is observed between the exposures at 1450 K and 1800 K for all composites. At 1250 K the value attributed to the sample H100 (1.1 mg cm−2 h−1) is close to those assigned to HfB2SiC composites oxidized at 1450 K (from 0.8 to 1.1 mg cm−2 h−1) but a significant increase is noticed at 1550 K (6.1 mg cm−2 h-1) and 1800 K (20.4 mg cm−2 h-1). Assuming negligible change in mass gain rate during cooling, the values are in agreement with those presented in the literature [14]. Study of the oxidation kinetics by isothermal TGA with direct introduction of all compositions at 1450 and 1800 K are shown in Fig. 6a) and b) respectively. These kinetics are compared to the most oxidation resistant composition in the ZrB2-SiC system which is the composition Z70S30, containing 70 vol% ZrB2 and 30 vol% SiC. The comparison of all the HfB2-based materials shows the interest of using HfB2 instead of ZrB2. Indeed, even the monolith H100 possesses a higher oxidation resistance than Z70S30 at 1450 and 1800 K. In any event, concerning HfB2-based materials, all compositions exhibit a slowing down of the oxidation rate. Moreover, a decrease of the final mass gain is observed when the amount of SiC increases. Oxidation kinetics show only mass gains at 1450 K whereas mass losses are observed during the isothermal runs at 1800 K for the H85S15, H75S25 and H70S30 samples. Overall values of specific mass at the end of TGA isothermal runs at 1450 and 1800 K are very close to those obtained after the oxidation in the solar furnace at the same temperatures. Different oxidation behaviors depending on the temperature and the composition can be seen in these plots. According to Fig. 6b) and because of their similar oxidation kinetics, the H95S5 and H90S10 samples can be associated. For the same reason, the H85S15, H75S25 and H70S30 samples can be considered as similar whereas the H80S20 sample exhibits a single oxidation behavior with no evolution of mass at the end of the isothermal run. This constant trend, never highlighted in the literature as far as the authors are aware, can be easily associated with a high protection against oxidation. The three different oxidation behaviors, represented by the H90S10, H80S20 and H70S30 samples are plotted in Fig. 7. A previous work [3] highlighted the presence of a protective borosilicate glass, on the surface of oxidized HfB2-SiC composites, which limits the diffusion of oxygen into the bulk material. In this work, oxidized surfaces were examined by SEM (Fig. 8) and EDS (not shown here). At 1450 K, all surfaces consist of HfO2(s) in white and SiC grains in dark. Morphologies and sizes depend on the initial composition and  1849  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 3. SEM images using backscattered electrons of the surfaces of sintered bodies: H90S10 a), H80S20 b) and H70S30 c). Higher magnification on H70S30 using secondary electrons d) showing the localization of pores (arrows) and EDS spectra of: HfB2 1), SiC 2), HfC 3) and HfO2 4).  the oxidation temperature. When the temperature increases up to 1800 K the morphology of grains changes with the appearance of blisters of a glassy phase. Moreover, the number of blisters at the surface depends on the initial SiC content.  These results are in a good agreement with X-ray diffraction analyses. Patterns of the oxidized surfaces at 1450 K and 1800 K of the H80S20 sample are shown in Fig. 1d) and e) and compared to the corresponding sintered material. After oxidation, samples exhibit  Fig. 4. Evolution as a function of SiC content of: sintering temperature a) and diboride grain size compared to the milled HfB2 powder b).  1850  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 5. Evolution of  the mass gain rate as a function of the oxidation temperature and the composition.  monoclinic HfO2 peaks, but also HfB2 and HfC lines. It has to be noticed that the oxidation temperature enhances the intensities of peaks for HfO2 while it decreases the peak intensities for HfB2 and HfC. Because it is supposed to be amorphous, the glassy phase did not reveal any peak by XRD. Raman maps have been used to characterize the oxidized surface of the H80S20 sample at 1800 K in order to investigate the spatial localization of the main phases. But above all, the aim was to identify a  signature for glassy domains which has been described in the literature but not seen by the classical methods used in this work. As it was previously explained, a multivariate method has been chosen for data analysis. In front of the complexity and the inhomogeneity of the materials, the PCA method has been selected using the Wire 4.3 software data analysis. The PCA method rotates the data matrix onto a new coordinate system where the greatest variance by any projection of the data comes to lie on the first coordinate (1st PC), the second greatest  Fig. 6. Specific mass change of all samples oxidized during a 20 min isothermal run at: 1450 K a) and 1800 K b).  1851  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  phase can be deduced. However, this silica is not pure because of the absence of the Q4 mode. Therefore, another chemical element must be incorporated in the silica-based network, such as boron. In-depth oxygen diffusion was investigated by SEM and EDS through the observation of cross-sections of oxidized materials at 1800 K (Fig. 10). The thickness of the oxidized layer depends on the composition. The highest obtained value was for the H90S10 sample with a 10 µm oxygen diffusion layer. H80S20 presents a thickness of 8.5 µm and only 5 µm for H70S30. EDS analyses confirm that at 1800 K the oxide layer consists of porous HfO2 with probably some trapped gaseous species (B2O3(g) and CO(g)), and Si which was not detected.  4. Discussion  4.1. Sintering  The optimization of sintering parameters of HfB2-SiC composites has been carried out in accordance with the results presented in Fig. 4a) showing a decrease of sintering temperature with the increase of the silicon carbide content. In order to prevent any grain growth, the imposed temperatures during Spark Plasma Sintering were lowered with the addition of SiC as shown in Table 1. This behavior is consistent with the results obtained by Sciti et al. [24] for the sintering of (Hf/Zr)B2 + MoSi2. It can be explained by the fact that when the silicon carbide content increases, a greater connectivity between these particles and the diboride can be expected leading to a more effective removal of oxygen impurities and thus to a lower sintering temperature, observed in Fig. 4a). On the contrary, if the amount of SiC is not high enough, all the oxides will not be removed thus promoting coarsening of the diboride as shown by Fig. 4b). These observations show a direct correlation between the sintering temperature and the diboride grain growth. In view of the high density (> 99%) of all materials, the homogeneous distribution of phases and the low grain growth during spark plasma sintering, all samples can be reliably compared in terms of oxidation resistance.  Fig. 7. Specific mass change for the samples H90S10, H80S20 and H70S30 oxidized at 1800 K during a 20 min isothermal run.  variance, of that remaining, on the second coordinate (orthogonal to the first), and so on. The PCA method is used to show variance percent and the associated loadings do not directly represent chemical information but the “weight” of each wavenumber component by component. The data treatment revealed at least 8 different Principle Components (PC). Nevertheless, only three of them correspond to three different phases and the other ones to these same phases but in stressed areas in the material. Indeed, all the PC extracted from the analysis will not be presented, but only the component related to the phases of interest. Three maps are shown in Fig. 9. The first one (PC1) is unambiguously linked to HfO2 and is unsurprisingly the first component with a variance of 51.40%. This means that this is the main compound present at the surface of the sample. SiC can also be found but only with 3.36% of variance as the forth component (PC4), which means that this phase exists but in a small proportion compared to HfO2. The third one (PC8) which appears only as the eighth component with a variance percent of 0.67 can be attributed to the Si-O vibration. According to Manara et al. [23], three modes can be proposed, located at 956, 1040 and 1085 cm−1 and attributed to Q2, Q0 and Q3, respectively. Based on this analysis and EDS information, the presence of a silica-rich glassy  Fig. 8. SEM images using secondary electrons of the oxidized surfaces of samples: H90S10 a) and d), H80S20 b) and e) and H70S30 c) and f) at 1450 K and 1800 K, respectively.  1852  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 9. Raman mapping obtained on the surface of the H80S20 sample oxidized at 1800 K showing: obtained maps of the three components of interest PC1, PC4 and PC8 a) and the attributed loadings b).  1853  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  Fig. 10. SEM images using secondary electrons of cross-sections of oxidized samples at 1800 K: H90S10 a), H80S20 b) and H70S30 c). Elemental composition of oxygen d), hafnium e) and silicon f) of the sample H90S10.  4.2. Oxidation  Fig. 5 shows the importance of adding SiC to HfB2 based materials to improve the oxidation resistance at high temperature. Because this plot shows only the mass gain rate after oxidation, kinetics were followed by thermogravimetric analysis with insertion of samples directly into the hot furnace and holding for the same 20 min period (Figs. 6 and 7). It can be seen from the plots at 1450 K in Fig. 6a) that materials are resistant to oxidation, more specifically the H80S20, H75S25 and H70S30 samples which exhibit almost no mass evolution after around 7 min exposure. These isothermal runs can be divided into three zones representing different regimes. At 1450 K and according to Eqs. (2) and (3) from the start up to 1 min, there is a reaction regime with formation of HfO2(s), SiO2(l) at the surface and volatilization of B2O3(g) and CO(g). Then the second zone corresponds to a complex regime with reaction and diffusion. Indeed, once the first oxidized condensed phases are formed, molecular oxygen can diffuse through them. Therefore, a part of oxygen will react with the non-oxidized grains at the surface and the other part will diffuse through hafnia pores or the silica layer. The last regime depends on the composition. For the H100, H95S5, H90S10 and H85S15 samples, a parabolic regime corresponding to the oxygen diffusion through the oxide layer can be observed whereas the H80S20, H75S25 and H70S30 samples do not show any mass evolution. In this work, this specific trend is called the constant regime. In brief, at 1450 K the initial amount of SiC governs the oxidation resistance. More specifically, with 20, 25 and 30 vol% SiC, materials present a good oxidation resistance with no evolution of mass. Similarly to Fig. 6a), three zones can be distinguished at 1800 K (Fig. 7). From the start up to 1 min of oxidation, a reaction regime takes place (I) and then a complex evolution occurs corresponding to a reaction and diffusion regime (II). The last zone (III) depends on the composition. Indeed, the H90S10 sample exhibits a diffusion regime with a mass gain whereas the H70S30 sample displays a linear evolution of mass loss from 5 min of oxidation called the recession regime. While this last composition presents a high oxidation resistance at 1450 K, this sample is not protective at 1800 K anymore because of the mass loss. Moreover, similar to its behavior at 1450 K, H80S20 does not  show any mass evolution at 1800 K what makes this composite the most resistant one to oxidation. Two hypotheses can thus be proposed here for this specific evolution. The first one is that the oxygen diffusion is stopped during the constant regime and thus there is no evolution of mass. The second one is that there is a mass compensation between the volatilization of gaseous species and the formation of condensed phases. The characterization of the surfaces and cross-sections may help in understanding the associated mechanisms. Figs. 8 and 9 show that the H80S20 surface oxidized at 1800 K during 20 min is composed of crystalline HfO2, and blisters made of different structures of a modified silica glassy phase. It has to be noted that the well-known Si-O-B band, characteristic of the borosilicate glass and located at 935 cm−1 [25], did not appear in this work. Based on these results, the presence of boron within the vitreous network, linked to Si-O bonds can not be confirmed. In addition, Raman spectroscopy did not highlight the presence of B2O3(l) in the H80S20 sample oxidized at 1800 K (B-O bands at 720 and 1265 cm−1) [26]. The absence of the B-O bands could reveal the volatilization of B2O3(g) below 1800 K. Furthermore, after a 20 min exposure at 1800 K, EDS analyses carried out in cross-section of H80S20 reveal the existence of a 8.5 µm thick SiC depletion layer, constituted of porous HfO2. This layer highlights the consumption of the composite by oxygen and the vaporization of SiO(g). Therefore, the second hypothesis which proposes a compensation between the volatilization of gaseous products (B2O3(g), CO(g) and SiO(g)) and the formation of condensed phases (HfO2(s) and SiO2(l)) can be validated. As a consequence, the modified silica glassy phase at the surface is not able to stop the oxygen diffusion. Nevertheless, the existence of a constant regime demonstrates a high protection against oxidation of the H80S20 sample up to 1800 K. Indeed, the fact that there is no evolution of mass at all for this composite would allow its reuse many times. The highest final mass was obtained for the H90S10 sample and can be correlated to the highest oxide thickness of 10 µm (Fig. 10). This expresses the less resistant nature of this composition with the possibility for oxygen to diffuse in depth up to 10 µm. Due to its lower initial amount of SiC and the absence of burst bubbles at the surface, an oxidized surface of the H90S10 sample should exhibit less blisters at the surface. This hypothesis is confirmed by Fig. 8 and is in a good  1854  \\x0c', 'C. Piriou et al.  Ceramics International 45 (2019) 1846-1856  agreement with Fig. 10. Indeed, the initial amount of SiC might not be high enough to create a highly protective surface but on the contrary, it should form a non-negligible amount of porous HfO2(s). Such a scale allows oxygen diffusion through interconnected pores [3] because the diffusion through the hafnia skeleton is negligible [27]. Therefore, because the oxygen transport is facilitated and the surface is not protective enough, it creates a thicker oxide layer. Nevertheless, the oxide thickness at this temperature is much lower than the values presented in the literature [12]. This discrepancy could be explained by a lower relative density obtained in their work (88-90%) compared to our study (> 99%). Here, fully-dense materials lead to a thinner oxide layer. The smallest oxygen layer thickness is obtained for the H70S30 sample. This composition containing the highest initial amount of SiC is subject to a recession regime at 1800 K that characterized species loss and thus a poor oxidation resistance. Moreover, after a 20 min exposure at 1800 K some burst bubble traces were detected on its oxidized surface. This could indicate the volatilization of gaseous species that probably drags in small hafnia grains during oxidation. Indeed, H70S30 possesses an initial SiC amount of 30 vol% contributing to a higher production of SiO2(l) and CO(g) than the two other compositions. As previously explained, the oxygen partial pressure in depth is much lower than at the surface which could imply an active oxidation of SiC and thus the volatilization of SiO(g) and CO(g) (Eq. (3)). During this phenomenon, gaseous species formed in depth will reach the surface and could cause a pressure increase on the external oxide. If the gas pressure exceeds the atmospheric pressure, bubbles made of glass will burst, exposing the underlying layer. If the bubble bursts during the oxidation, the surface could be re-oxidized but if it bursts during cooling, the surface will stay non-oxidized. Indeed, after the oxidation in the REHPTS facility, cooling is very fast which means that the temperature should be lower than the oxidation temperature of HfB2 and SiC. The volatilization phenomenon could counter oxygen flow, slow down oxygen diffusion through the oxide layer and thus it could slow down the formation of SiO(g) by active oxidation, explaining the smaller thickness of the oxide layer. It has to be noticed that viscosity of the glass also has an influence on bubbles bursting [28] but it is not discussed here. In brief, because of its low stability, the H70S30 composition cannot be used for the protection against oxidation at 1800 K.  5. Conclusion  Fully-dense and microstructure-controlled HfB2-SiC composites from 0 up to 30 vol% SiC were successfully achieved by Spark Plasma Sintering. A decrease of the sintering temperature with the amount of SiC was observed and directly connected to the removal of oxide impurities located at the surface of the diboride particles. In order to prevent any grain growth during the densification, sintering parameters were optimized by decreasing the imposed temperatures with the amount of SiC. No evidence of excessive grain growth was highlighted for the composites or the monolith. Silicon carbide was also established as being useful for the oxidation resistance with an overall decrease of the mass gain. All compositions exhibited a protection against oxidation at 1450 K and a higher resistance was observed when the amount of SiC increased. However, different oxidation behaviors were highlighted at 1800 K depending on the composition. The H80S20 sample turned out to be the most resistant one with no evolution of the mass gain after 7 min exposure. This high protection against oxidation is promoted by the presence of a covering of modified silica glassy phase at the surface which limits the oxygen diffusion in depth. Indeed, the formation of a such protective external oxide generates longer diffusion paths for oxygen to reach the boundary between the oxide and the non-oxidized composite [7]. The study of the oxidation mechanisms shows that this constant regime could be due to a compensation between the volatilization of gaseous species (B2O3(g), CO(g) and SiO(g)) and the formation of condensed phases (HfO2(s) and SiO2(l)). The H90S10 sample oxidized at 1800 K, for its part, exhibited a parabolic regime after 11 min  1855  oxidation and possesses the thickest oxide layer with a 10 µm oxygen diffusion zone. Moreover, even if the 30 vol% SiC containing composite was the most resistant one at 1450 K, it presented a mass loss after 5 min exposure at 1800 K indicating a poor oxidation resistance at this higher temperature. From all these results, it can be confirmed that the composition has a real influence on the oxidation behavior. The oxidation behavior of these compositions, and more specifically of H80S20, could be investigated at higher temperatures and during several cycles in order to highlight their maximum abilities under oxidation.  Acknowledgments  The authors wish to thank the GDR CNRS No. 3584 TherMatHT for fruitful discussions and collaborative work on the present project. They are also grateful to Julie Cornette for her help and advice in Raman spectroscopy.  References  [7]  [2]  [3]  [5]  [6]  [10]  [1] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. Soc. 19 (1999) 2405-2414. E.V. Clougherty, K.E. Wilkes, R.P. Tye, Research and development of refractory oxidation-resistant diborides. Part II. Volume V. Thermal, physical, electrical, and optical properties, MANLABS, INC., CAMBRIDGE, MASS., AD-−865321; AFML-TR−68-190, 1969. F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC composite, J. Eur. Ceram. Soc. 25 (2005) 1025-1031. [4] X. Ren, H. Li, Y. Chu, Q. Fu, K. 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},{
  "_id": 244,
  "PDF": "Solid‐Solution Effects on the High‐Temperature Oxidation Behavior of Polymer‐Derived (Hf,Ta)C-SiC and (Hf,Ti)C-SiC Ceramic Nanocomposites.pdf",
  "Text": "['FULL PAPER  Advanced Non-Oxide Ceramic Matrix (Nano)-Composites  www.aem-journal.com  Solid-Solution Effects on the High-Temperature Oxidation Behavior of Polymer-Derived (Hf,Ta)C/SiC and (Hf,Ti)C/SiC Ceramic Nanocomposites  Qingbo Wen, Ralf Riedel, and Emanuel  Ionescu*  In the present study,  two concepts to improve the oxidation resistance at  high-temperatures of ceramic nanocomposites consisting of 85-90 vol% SiC,  5-8 vol% group IV metal carbides (i.e., HfC, TaC), and 5-7 vol% carbon are  introduced and discussed. First  improvement concept relates to the passiv ation of  the samples upon short-term oxidation at 1400  C (30 min). This is a  \\x0e  critical step, especially with respect  to silica formation, which is relatively  sluggish at  temperatures lower  than 1000-1200  C. Moreover, solid-solution  \\x0e  metal carbides (Hf,Ta)C and (Hf,Ti)C are shown to be clearly more oxidation  \\x0e  \\x0e  However, poor oxidation resistance as well as bad damage tolerance of HfCbased UHT ceramics and related materials limit their practical applications.[13-15] In order to improve their environmental stability, within the context of (ultra) high-temperature applications, UHTCbased formulations such as UHTC (nano)composites[9,16-25] and CMCs[26,27]) were developed. Typically, a secondary silica-forming phase is added to the UHTC, which is expected to protect the UHTC phases by forming a dense Sibased scale with low oxygen diffusivity at intermediate exposure temperatures (1000-1600 C).[24,26-44] The Si-based protective scale (e.g., SiO2, borosilicate, metal silicates) is supposed to protect the UHTCs at temperatures up to 1600 C; whereas beyond this temperature the UHTC phase is expected to deliver oxide scales being able to densify fast enough and consequently provide environmental robustness at ultrahigh temperatures.[27,45-49] An additional effect of using secondary silica-former phases relates to the reduction of the density of the corresponding UHTC (nano)composites.[50,51] In our recent work, dense monolithic (Hf,Ta)C/SiC-based ceramic nanocomposites were synthesized and investigated concerning their high-temperature oxidation resistance.[12,52] It was shown that the oxidation behavior of HfC/SiC at temperatures from 1200 to 1500 C is rather poor though it can be significantly improved upon alloying of HfC with TaC, probably due to the significantly higher densification ability of the oxides generated upon oxidation of the metal carbide phases in (Hf,Ta) C/SiC nanocomposites, which are able to provide together with and hinder oxidation.[52] silica a protective continuous scale However, it was observed that all studied nanocomposites are prone to significant oxidation accompanied by the generation of porosity and cracks especially at temperatures below 1200 C and was concluded that in this temperature range the formation of silica from the SiC phase is not activated enough to provide a continuous protective scale.[52] In order to solve this problem, further investigation on oxidation behavior of polymer-derived (Hf,Ta)C/SiCand (Hf,Ti)C/SiC-based ceramic nanocomposites has been conducted in the present work. We propose and critically assess a synergistic strategy of combining (M ¼ Ta, Ti) with a (Hf,M) solid-solution effect passivation procedure of the samples prior to the exposure to the high \\x0e  \\x0e  resistant  than the binary HfC and TaC phases. Whereas,  the solid-solution  effect contributes to a significant  improvement of  the short-term oxidation  resistance of  the studied nanocomposites,  the passivation of  the materials  prior exposure of high-temperature oxidation conditions provides a remark ably improved long-term behavior thereof. Possible mechanisms involved in  the oxidation processes of  (Hf,Ta)C/SiC and (Hf,Ti)/SiC ceramic nano composites are highlighted and critically assessed.  1.  Introduction  Hafnium carbide and its solid solutions with TaC and TiC [i.e., (Hf,Ta)C and (Hf,Ti)C] are well known advanced ceramics which attract much attention in the past decade.[1-5] Particularly, the physical and chemical properties of the ternary carbides vary as a function of the molar ratio of the metal elements (i.e., Hf to Ta and Hf to Ti) and are always better than those of the binary constituents (i.e., HfC, TaC, and TiC).[6] For instance, Ta4HfC5, that is, (Hf0.2Ta0.8)C, has been proven to possess the highest melting point (4215 K) among all currently known materials.[7,8] A hardness of approximately 43 GPa is reported for (Hf0.5Ti0.5)C, making it one of the hardest materials known.[1] Moreover, HfC, TaC, TiC, and their solid solutions exhibit excellent electrical \\x001, which makes it possible conductivity in the order of 104 S cm to be used as conductive ceramics or electromagnetic wave absorbing/shielding materials in harsh environment (i.e., in corrosive/oxidizing media or at high temperatures).[1,9-12]  Dr. E. Ionescu, Dr. Q. Wen, Prof R. Riedel Technical University of Darmstadt Institute of Materials Science Otto-Berndt-Str. 3, Darmstadt D-64287, Germany E-mail: ionescu@materials.tu-darmstadt.de  DOI: 10.1002/adem.201800879  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (1 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  temperature oxidation conditions in order to substantially enhance the long-term oxidation resistance of the prepared ceramic nanocomposites over a broad temperature range up to 1500 C.  \\x0e  2. Experimental Section  2.1. Materials Synthesis and Processing  [9,12].  Dense monolithic (Hf,Ta)C/SiCand (Hf,Ti)C/SiC-based ceramic nanocomposites were prepared upon spark plasma sintering of amorphous Six(Hf,Ta)yCz and Six(Hf,Ti)yCz ceramic powders that were obtained via pyrolysis of (Hf,Ta)and (Hf,Ti)containing preceramic single-source precursors, respectively.[9,12] Following ceramic nanocomposites were prepared and investigated in the present study: TaC/SiC, (Hf0.2Ta0.8)C/SiC, (Hf0.7Ta0.3)C/SiC, HfC/SiC, (Hf0.9Ti0.1)C/SiC, and (Hf0.5Ti0.5)C/ SiC. The chemical synthesis of the Hfand (Hf,Ta)-containing single-source precursors was previously reported in refs. Similarly, the (Hf,Ti)-containing single-source precursors were synthesized upon reacting a commercially available allylhydridopolycarbosilane (SMP10, Starfire System Inc, USA) with amido complexes of Hf and Ti (i.e., Hf(NEt2)4 and Ti(NEt2)4). The weight ratio of pure Hf(NEt2)4 complex to SMP10 was set to be 30:70, and the molar ratio of Hf(NEt2)4 to Ti(NEt2)4 for preparation of the (Hf0.9Ti0.1)C/SiC and (Hf0.5Ti0.5)C/SiC ceramic nanocomposites was set to be 9:1 and 1:1, respectively.[9,53] All Hf-containing monoliths were sintered in vacuum at 2200 C and 50 MPa for 20 min with a FCT HP D 25/1 SPS apparatus (FCT Systeme GmbH, Frankenblick, Germany). The \\x001, heating and cooling rates were 100 and 220 C min respectively. The TaC/SiC monoliths were sintered using the same condition using a SPS-1050 apparatus (SPS Syntex Inc., Kawasaki, Japan). The obtained ceramic monoliths to be used for oxidation tests were cut into specimens with the dimension of 4 \\x02 3 \\x02 2 mm and carefully polished with 1 mm polycrystalline diamonds on felt cloth. Before the oxidation test, the specimens were dried at 80 C for 48 h and then weighed on a semi-micro balance (0.01 mg).  \\x0e  \\x0e  \\x0e  2.2. Materials Characterization  Elemental analysis, X-ray diffraction (XRD) and scanning electron microscopy (SEM) coupled with EDX analysis were used to characterize the phase composition, surface morphology and microstructure of the specimens prior and after oxidation. The carbon contents of the ceramics were measured using a combustion analysis method on a LECO C-200 analyzer (LECO Corporation, St. Joseph, Michigan, USA). The nitrogen and oxygen contents were measured with a LECO TC-436 analyzer (LECO Corporation, St. Joseph, Michigan, USA). Hf, Ti, and Si content were measured using Inductively Coupled PlasmaAtomic Emission Spectrometry (ICP-AES) on an iCap 6500 device (Thermo Fisher Scientific, Waltham, Massachusetts, USA). The XRD patterns of as-synthesized ceramic powders  were recorded using a STADI P powder diffractometer (STOE & Cie GmbH, Darmstadt, Germany) with a molybdenum Kα1 (λ ¼ 0.709300 A radiation source ). The XRD patterns of the monolithic samples materials were measured by a Bruker D8 system (Bruker Corporation, Billerica, Massachusetts, United States) with a copper Kα1 radiation source (λ ¼ 1.541874 A ). The surface morphology and fractured cross section of the specimens were investigated by a Philips XL30 FEG high-resolution scanning electron microscope (FEI Company, Hillsboro, Oregon, USA) coupled with an Energy-dispersive X-ray spectroscope (EDAX, Mahwah, New Jersey, USA). Measurements of the skeletal density and open porosity of the prepared bulk materials have been performed using the water immersion (md) of the sample pellets have method. Thus, size and weight been first measured. Subsequently, the samples were heated in boiling water for 60 min in order to fill the open pores and the the wet samples (mw) and the weight of weight of the samples immersed in water (mimm) have been determined. The skeletal densities (ρs) and open porosity (p) values of the materials have been calculated using the following equations:  ρ ¼  md  md \\x00 mimm  g cm3  p %ð  Þ ¼  mw \\x00 md  mw \\x00 mimm  \\x02 100  ð1Þ  ð2Þ  2.3. High-Temperature Oxidation Experiments  In situ assessment of the high-temperature oxidation of selected samples was performed on a thermogravimetric analysis (TGA) (STA 449C Jupiter, Netzsch Gera ̈tebau GmbH, Selb/ device \\x001). The samples Bavaria, Germany) were heated from \\x1920 in dry air flow (75 mL min \\x001 and subsequently held at C to 1400 or 1500 C with the heating rate of 5 C min the target temperature for 5 h.  \\x0e  \\x0e  \\x0e  2.4. Surface Passivation of  the Monolithic Samples  \\x0e  The passivation process for all monolithic specimens was conducted at 1400 C for 5 h in an alumina furnace. A SiC crucible was used as a container to minimize impurity effects. The specimens were placed in the isothermal area of the furnace when the target temperature was achieved and subsequently held for 30 min. The weight of the specimens after pre-oxidation was recorded on a semi-micro balance (0.01 mg). The surfacepassivated monoliths are denoted as p-TaC/SiC, p-(Hf0.2Ta0.8)C/ SiC, p-(Hf0.7Ta0.3)C/SiC, p-HfC/SiC, p-(Hf0.9Ti0.1)C/SiC, and p(Hf0.5Ti0.5)C/SiC.  2.5.  Isothermal Oxidation of  the Passivated Samples  The experiments were conducted in air at 1200, 1300, 1400 and 1500 C, respectively. The weight of the specimens was  \\x0e  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (2 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  ̊ ̊ \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  measured after oxidation times of 5, 15, and 20 h. Moreover, isothermal oxidation tests for 100 h were conducted at 1400 C. The apparent parabolic rate constant (Kp) of the isothermal oxidation test was calculated using Equation (3).  \\x0e  \\x12  \\x13  K p \\x03t ¼  ΔW  A  2  ð3Þ  where, ΔW=A is the mass change per unit of surface area in mg \\x002, A the cm surface area (cm2), and t the oxidation time in hours.[54] The surface area (A) was considered to be the geometric surface area (cm2) of the monoliths. Using the parabolic rate constants at different temperatures, the activation energy (Ea) of the oxidation process was calculated as  K p ¼ A0 e\\x00Ea  RT  ð4Þ  where R is the ideal gas constant absolute temperature (K).  (\\x19 8.314 J/(mol\\x01K)) and T the  3. Results and Discussion  3.1. Structural Characterization of  the Studied Materials  XRD patterns of the as-sintered monoliths used in the present study are shown in Figure 1. The XRD results indicate that all sample consist of SiC and metal carbide phases, in addition to small amounts of excess carbon. The reflections of the metal carbide phase appear at 2θ values which correlate to the extent of modification of the HfC phase with Ta and Ti. Thus, the lattice parameter of HfC (4.63039 A , see Table 1) changes upon modification with 30 and 80 mol% Ta to 4.5792 and 4.4943 A as expected. This trend can be rationalized with the Vegard’s rule.[55] The open porosity, average grain sizes and phase composition of the monoliths used for oxidation test are shown in Table 1. All monoliths are nearly dense, as they exhibit residual open porosities less than 1.5 vol%. The average grain sizes and lattice parameters of the (Hf, Ta) (Hf, Ti)C, and β-SiC phases were obtained by Full-Profile C, Rietveld refinement of the XRD patterns of the investigated samples. The peak shapes were modeled using the ThompsonCox-Hastings pseudo-Voigt function for the average grain sizes and using the pseudo-Voigt function for lattice parameters. The phase composition (volume percent) was estimated using the results of elemental analysis and the density of the (Hf, Ta)C, (Hf, Ti)C, and β-SiC phases were calculated also by Rietveld \\x003 was used as the refinement. A density of graphite of 2 g cm density of the free carbon in the studied materials. The values of densities and mass fractions of the various phases in the prepared nanocomposites are summarized in Table 2. The estimated average grain sizes of both (Hf,Ta)C and (Hf,Ti) C grains within all the monoliths are less than 100 nm (Table 1). The phase composition of the monoliths consisted of β-SiC as the main phase (i.e., 85-89 vol%) as well as (Hf,Ta)C or (Hf,Ti)C (5-7 vol%) and segregated carbon (6-8 vol%). It can be concluded that the single-source precursor synthesis approach  Figure 1.  SiC [9,12]  a) XRD patterns of the as-sintered dense monolithic (Hf,Ta)C/  and  (Hf,Ti)C/SiC-based  ceramic  nanocomposites  used  for  oxidation test; b) magnification of  the XRD patterns showing the (111)  and (200) reflections of  the (Hf,Ta)C phases.  \\x0e  used in this study provides a facile access to nanocomposites based of SiC and solid solutions of HfC with TaC/TiC. The solidsolution metal carbides do not show any phase separation during the SPS densification processes conducted at 2200 C. Interestingly, the reflections of the metal carbide phases in the prepared nanocomposites show significantly higher intensities as compared to those of β-SiC, despite the weight fractions of the (Hf,M)C phases are relatively low (< 25 wt%). Indeed, one cannot simply compare the weight fraction of different phases simply by comparing the intensity of their reflections; however, it seems to be rather contra-intuitive. Nevertheless, Rietveld refinement of the XRD patterns were in agreement with the elemental analysis data and indicated the phase composition shown in Table 1. The fact that the reflection intensity of the (Hf, M)C phases is higher than that of that of β-SiC can be explained by the fact that the (Hf, M)C phase, except for (Hf0.5Ti0.5)C, is significantly better crystallized than β-SiC. In the case of (Hf0.5Ti0.5)C/SiC, we assume that the crystallization of  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (3 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  ̊ ̊ \\x0c', '(Hf0.5Ti0.5)C is suppressed due to the presence of large amounts of Ti.  3.2. TGA-Assisted Oxidation Tests of  the As-Prepared  Ceramic Monoliths  Figure 2 shows the oxidation behavior of HfC/SiC and TaC/SiC monoliths (oxidation curves in Figure 2a and parabolic plots in Figure 2b), which was studied via thermogravimetric analysis in air. As shown in the non-isothermal part of the oxidation curve of HfC/SiC (Figure 2a), there is a mass gain process occurring in the temperature range from ca. 530 C to ca. 700 C. This correlates to the oxidation of HfC, which occurs at relatively low temperatures (500-600 C),[56] and is rather fast at the beginning, leading to a significant mass gain which is subsequently counteracted by the rapid oxidation of the sp2 carbon phase present in the material (temperature range from 700 to ca. 900 C, Figure 2a). For TaC/SiC, the oxidation curve shows mass loss process(es) in the temperature range from 800 to 1200 C, which probably relies on the oxidation of sp2 carbon, and adopts a parabolic evolution (mass gain) at higher temperatures. The oxidation of  \\x0e  \\x0e  \\x0e  \\x0e  \\x0e  TaC at low temperatures is not visible in the oxidation curve of TaC/SiC. The apparent parabolic rate constant Kp of HfC/SiC (3.69 mg2/(cm4 h), Figure 2b) indicates that its oxidation kinetics is rather fast as compared to that or sintered SiC[59] exposed to oxidation condition at the same \\x005-10 temperatures (with reported Kp values in the order of 10 mg2/(cm4 h)). It should be noted that HfC/SiC consists mainly of SiC (ca. 87.5 vol%, see Table 1); however, the presence of HfC (ca. 7 vol%) and sp2 carbon (ca. 5.5 vol%) significantly accelerates the oxidation in HfC/SiC as compared to SiC. The Kp value of TaC/ \\x002 mg2/(cm4 h), Figure 2b) was SiC (2.7 10 found to be ca. 2 orders of magnitude lower than that of HfC/SiC, despite showing a similar phase composition (ca. 89 vol% SiC, 5 vol% TaC, and 6 vol% sp2 carbon). This significant difference has been attributed to the strongly different melting temperatures and consequently different densification behavior of HfO2 and Ta2O5 in the scale of HfC/SiC and TaC/SiC, respectively, at the investigated temperatures. Thus, Ta2O5 has a significantly lower melting temperature (ca. 1872 C) than that of HfO2 (ca. 2758 C) and is consequently able to densify faster. However, TaC/SiC also shows apparent parabolic oxidation rates being 2-3 orders of magnitude higher than those of dense SiC.  of CVD SiC[28,54,57,58]  \\x003  \\x0e  \\x0e  Table 1. Open porosity, average grain sizes,  lattice parameter as well as elemental and phase compositions of  the as-prepared ceramics.[9,12]]  Average grain size [nm]a)  Phase Composition [vol%]b)  Sample  Open porosity  [vol%]  (Hf,Ta)C/(Hf,Ti)  C  SiC  Lattice parameter of  (Hf,Ta)C [A  ]  Empirical  formulae  SiC  (Hf,Ta)C/(Hf,Ti)C  Free C  TaC/SiC  \\x190  73.7  75.3  4.4584  SiTa0.054C1.195(O0.131)  89.0  4.9  6.1  (Hf0.2Ta0.8)C/  SiC  1.14  72.3  73.0  4.4943  SiHf0.010Ta0.043C1.243(O0.104)  87.5  4.9  7.7  (Hf0.7Ta0.3)C/  SiC  1.40  91.1  81.0  4.5792  SiHf0.041Ta0.018C1.240(O0.078)  86.8  5.8  7.4  HfC/SiC  0.74  80.9  36.6  4.6309  SiHf0.070C1.194(N0.016O0.125)  87.5  7.0  5.5  (Hf0.9Ti0.1)C/  SiC  1.05  32.4  71.1  4.6107  SiHf0.047Ti0.005C1.234(O0.124)  87.3  5.2  7.5  (Hf0.5Ti0.5)C/  SiC  0.16  13.2  105.5  4.5013  SiHf0.041Ti0.040C1.269(N0.012O0.116)  84.9  7.3  7.8  a) The average grain sizes are estimated by Rietveld refinement of the XRD patterns using Full-Profile software. b) The phase composition was estimated using the results of elemental analysis.  Table 2. Calculated densities (from Rietveld refinement) and mass fractions (from elemental analysis) of  the phases present  in the studied  nanocomposites.  SiC  (Ta,Hf )C/(Hf,Ti)C  Carbon  Sample  Density [g cm-3]  Mass fraction [wt%]  Density [g cm-3]  Mass fraction [wt%]  Density [g cm-3]  Mass fraction [wt%]  TaC/SiC  3.220  77.6  14.462  19.1  2  3.3  (Hf0.2Ta0.8)C/SiC  3.209  77.0  14.085  18.8  2  4.2  (Hf0.7Ta0.3)C/SiC  3.211  75.3  13.228  20.7  2  4  HfC/SiC  3.222  73.7  12.845  23.4  2  2.9  (Hf0.9Ti0.1)C/SiC  3.207  78.4  12.025  17.4  2  4.2  (Hf0.5Ti0.5)C/SiC  3.213  76.8  9.121  18.8  2  4.4  www.advancedsciencenews.com  www.aem-journal.com  Adv. Eng. Mater. 2019, 21, 1800879  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  1800879 (4 of 11)  ̊ \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  Figure 2.  a) Oxidation curves of  the as-prepared HfC/SiC and TaC/SiC  monoliths (grayish area represents the isothermal part of  the oxidation  test performed at 1400  C); b) Parabolic plot of the isothermal part of the  C of  the  as-prepared HfC/SiC and  TaC/SiC  \\x0e  \\x0e  oxidation  test  at  1400  monoliths.  \\x0e  \\x0e  As previously reported,[52] the (Hf,Ta)C/SiC-based ceramic nanocomposites exhibit excellent oxidation resistance at temperatures from 1300 to 1500 C but poor oxidation resistance at 1200 C due to the rapid oxidation of the (Hf,Ta)C phase and segregated carbon and limited sintering/densification ability of the resulting metal-containing oxides. Similar oxidation behavior can be observed for the (Hf,Ti)C/SiC-based ceramic nanocomposites (Figure 3). The parabolic rate constants reveal that the oxidation resistance of the (Hf,Ti)C/SiC monoliths at 1400 and 1500 C is much better than that at 1200 and 1300 C (Table 3). Particularly, the Kp values at 1200 C of all samples containing solid-solution \\x002 mg2/(cm4 h) and metal carbides are in the order of 100-10 some of the monoliths show significant cracking after oxidation (see Supporting Information, Figure S1).[52] The photographs and SEM micrographs of the oxidized HfC/ SiC and TaC/SiC (Figure 4) clearly show that both samples are seriously damaged upon oxidation at 1400 C. As shown in Figure 3a, the color of the oxidized HfC/SiC turned from black to gray upon oxidation and the sample catastrophically cracked after the test. TaC/SiC is still black, but one large crack can be  \\x0e  \\x0e  \\x0e  \\x0e  Figure 3. Oxidation curves of a)  (Hf0.9Ti0.1)C/SiC and b)  (Hf0.5Ti0.5)C/  SiC.  observed on the surface (Figure 4b, left). Moreover, the SEM micrograph of the cross section of TaC/SiC (Figure 4b, right) reveals pronounced cracking in the interior of the sample. The failure of the HfC/SiC sample can be understood as a consequence of the oxidation of HfC which delivers porous hafnia scale, unable to protect the monolith from further oxidation; whereas the pronounced cracking of the surface and  Table 3. Apparent parabolic oxidation rated Kp of monolithic samples.  the as-prepared  Sample  TaC/SiC  (Hf0.2Ta0.8)C/SiC  (Hf0.7Ta0.3)C/SiC  HfC/SiC  (Hf0.9Ti0.1)C/SiC  (Hf0.5Ti0.5)C/SiC  \\x005 mg2/(cm4 h)] Apparent parabolic oxidation rate Kp [10  1200  \\x0e  C  123  2800  3800  21 300  8500  10 300  1300  \\x0e  C  1400  \\x0e  C  1500  \\x0e  C  13.2  90.3  75.9  8.65  45.7  41.7  5.34  34.8  34.8  26 800  19 200  21 200  4300  350  45.6  11.7  29.8  5.8  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (5 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  Figure 4. Photographs of a) HfC/SiC and b) TaC/SiC after oxidation at 1400  \\x0e  C; c) SEM micrograph of the cross section of the oxidized TaC/SiC sample  shown in (b).  \\x0e  within the scale of TaC/SiC can be correlated to the intrinsic behavior of the Ta2O5 which is generated upon oxidation of TaC and shows a good match with SiC concerning the coefficient of thermal expansion but undergoes a phase transformation at high temperatures (i.e., orthorhombic ! tetragonal, i.e., β ! α, at ca. 1360 C) and thus suffers from big volume changes during the oxidation process.[60] In conclusion, HfC/SiC shows a poor oxidation behavior because the generated HfO2 is porous and not able to densify quickly at the temperatures used for the experiments. In the case of TaC/SiC, Ta2O5 may be able to densify fast (this is reflected in the parabolic oxidation rates being significantly lower than those of HfC/SiC), however it suffers from significant cracking occurring as a consequence of reversible phase transformations at high-temperatures. In the following, a combination of two tools/concepts to improve the oxidation behavior of HfC/SiC is introduced and discussed. One concept relates to the improvement of the oxidation resistance via “passivation” of the surface of HfC/SiC upon short-term oxidation at high temperatures. The strategy involves the activation of the oxidation process of the β-SiC phase and consequently the generation of a continuous scale during the passivation, which is the metal carbide and sp2 carbon phases against able to protect oxidation at low temperatures more effectively than just protecting those via encapsulation in the SiC matrix. The second concept generating solid-solution (Hf,M)C phases (M ¼ Ta, Ti). Here, involves the modification of the HfC phase in HfC/SiC upon it is shown that the metal used for the generation of the solid solution determines the mechanism(s) of improving the oxidation resistance of HfC/SiC, as discussed below.  3.3. Passivation of  Exposure to 1400  \\x0e  the Monoliths upon Short-Time  C  \\x0e  As presented above, HfC/SiC and TaC/SiC seriously cracked during the TGA-oxidation test due to the rapid oxidation of HfC, TaC, and segregated carbon in the temperature range from ca. 500 to 800 C. The β-SiC matrix can effectively form a protective SiO2 film only at temperatures higher than 1100 C. Thus, the oxidation processes of HfC, TaC, and segregated carbon occurring at moderate temperatures lead to various cracks and pores which provide fast track pathways for oxygen further accelerate the oxidation reaction.  \\x0e  \\x0e  \\x0e  \\x0e  XRD patterns of the passivated samples (i.e., exposed for 30 min to 1400 C in air) are presented in Figure 5. Sample pTaC/SiC shows the presence of β-Ta2O5 and small amounts of βcristobalite on its surface (Figure 5a, b). As stated above, the formation of β-Ta2O5 may be attractive due to its facility to quickly densify at temperatures of 1300-1500 C; however, it is detrimental due to its phase transformation to α-Ta2O5 at temperatures beyond 1360 C which is accompanied by large volume change (8 vol%). On the surface of the p-(Hf0.2Ta0.8)C/SiC and p-(Hf0.7Ta0.3)C/ SiC, the formation of Hf6Ta2O17 and TaC can be observed in addition to β-Ta2O5 and β-cristobalite (Figure 4a, b).[61] Following processes are assumed to occur during the oxidation of (Hf0.2Ta0.8) C/SiC and (Hf0.7Ta0.3)C/SiC: in a first step, the (Hf,Ta)C phase is \\x00 this can be seen as a result of a selective oxidation of Ta in (Hf,Ta) assumed to oxidize (Equation (5) and (6)) to furnish Ta2O5 and HfC solid solutions, leading to Ta-depleted HfC in addition to Ta2O5. Subsequently, HfC may react with Ta2O5 to furnish HfO2 and TaC (Equation (8)), as also reported previously.[62] This is the only reasonable explanation for the formation of TaC during a hightemperature process occurring in oxidative conditions. Indeed, the change of the free Gibbs energy for the Equation (8) was calculated with Factstage for the temperature range from 500 to 2000 C to be negative (cf. ΔG ¼-0.14 T-759.76 kJ) and thus thermodynamically favorable.[62] Additionally, HfO2 and Ta2O5 react to deliver Hf6Ta2O17 (Equation (7)). This rather complex evolution of the phase composition in the oxidation scale is considered to be beneficial for the oxidation behavior of (Hf0.2Ta0.8)C/SiC and (Hf0.7Ta0.3)C/SiC (as compared to that of TaC/SiC) due to following reasons: 1) the presence of Hf in the (Hf,Ta)C solid solution is responsible for the consumption of Ta2O5 during oxidation and thus acts against cracking of the scale due to phase transformation processes of Ta2O5; 2) formation of Hf6Ta2O17 is seen also as beneficial due to its high-temperature phase stability as well as its low thermal conductivity and oxygen diffusivity.[63-65] Moreover, it is shown that the composition of the scale can be controlled/tuned by the Hf:Ta ratio in the (Hf,Ta)C solid solution phase.  \\x0e  ðHf x Ta1\\x00x ÞC þ 7 \\x00 5x 2  O2 ! 1 \\x00 x 2  Ta2O5 þ xHfC þ CO  HfC þ 3 2  O2 ! HfO2 þ CO  ð5Þ  ð6Þ  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (6 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  Figure 5. XRD patterns of the (Hf,Ta)C/SiCa), b), and (Hf,Ti)C/SiC-based c) ceramic nanocomposites after passivation upon oxidation at 1400  \\x0e  5 h. The pattern shown in (b) is a magnification from (a) in the 2theta range from ca. 18 to 45  .  \\x0e  C for  6HfO2 þ Ta2O5 ! Hf 6Ta2O17  3Ta2O5 þ 7HfC ! 7HfO2 þ 6TaC þ CO  2SiC þ 3O2 ! 2SiO2 þ 2CO  ð7Þ  ð8Þ  ð9Þ  \\x0e  As an additional consequence of the passivation process performed at 1400 C, the formation of silica (cf. Equation (9)) is seen as beneficial for the oxidation behavior in the studied nanocomposites as it provides the matrix of the scale which is present as continuous phase and thus prevents from further oxidation of the nanocomposites. The XRD patterns of p-(Hf0.9Ti0.1)C/SiC and p-(Hf0.5Ti0.5)C/ SiC show the presence of hafnon (HfSiO4) in addition to m-HfO2 amounts of β-cristobalite and small (Figure 5c). The HfSiO4 phase has a melting point of ca. 1750 and it has a significantly lower oxygen diffusivity than that of silica.[68] Thus, the HfSiO4 phase can sinter effectively at the temperatures used  C,[66,67]  \\x0e  \\x0e  C.[69,70]  in the present study and thus may forms an excellent protective scale against oxidation together with SiO2. This is different from the situation in p-HfC/SiC, which does not show any formation of hafnon (only m-HfO2 is formed upon oxidation). Usually, the formation of HfSiO4 from HfO2 and SiO2 (similar to ZrSiO4 formation from ZrO2 and SiO2) is rather sluggish and not activated at temperatures below 1400 Here, the beneficial effect of the presence of Ti in the (Hf,Ti)C solidsolution phase is obvious, as it clearly facilitates the formation of hafnon, probably via an intermediary formation of glassy TixSiyOz which may react with HfO2 to HfSiO4.[70,71] Additionally, the formation of HfO2-TiO2 solid solutions via oxidation of (Hf,Ti)C and their subsequent reaction with SiO2 may be an alternative path to hafnon. Also, the Hf:Ti ratio seems to determine the phase composition of the scale. For p-(Hf0.9Ti0.1)C/SiC, small amounts in addition to m-HfO2 of HfSiO4 are detected and little cristobalite; whereas in p-(Hf0.5Ti0.5)C/SiC, HfSiO4 fraction is significantly increased and additionally the formation of HfTiO4 is observed. Based on the pseudo binary HfO2-TiO2 phase diagram, the solubility of TiO2 in m-HfO2 amounts ca. 10 mol%. For HfO2-TiO2 systems with higher TiO2 molar fractions, that is,  β Adv. Eng. Mater. 2019, 21, 1800879  1800879 (7 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  \\x0e  C/SiC due to following reasons: 1) it has a melting point of ca. 1980 C and consequently, it is expected that it effectively sinters at the temperatures used for the oxidation experiments; 2) it does not undergo any phase transformation up to its melting point; 3) it is thermoshock resistant (similarly to ZrTiO4), due to its high structural anisotropy.[73-75]  3.4. Short-Term Oxidation Behavior of Passivated Monoliths  at High Temperatures  \\x0e  \\x0e  \\x0e  The oxidation resistance of the passivated monoliths, including p-TaC/SiC, p-(Hf0.2Ta0.8)C/SiC, p-(Hf0.7Ta0.3)C/SiC, p(Hf0.9Ti0.1)C/SiC, and p-(Hf0.5Ti0.5)C/SiC, was tested using thermogravimetric analysis from 20 to 1500 C in dry air flow. The TG oxidation curves are shown in Figure 6. Compared with Figure 2a, no significant mass change has been detected in both non-isothermal and isothermal part of the oxidation tests. This leads to apparent parabolic rate constants (Kp) at 1500 C having values of 3.21 \\x02 10 \\x004, 4.52 \\x02 10 \\x004, 6.49 \\x02 10 \\x004, 2.11 \\x02 10 \\x003, and 6.50 \\x02 10 \\x004 mg2/(cm4 h) for p-TaC/SiC, p-(Hf0.2Ta0.8)C/SiC, p(Hf0.7Ta0.3)C/SiC, p-(Hf0.9Ti0.1)C/SiC, and p-(Hf0.5Ti0.5)C/SiC, respectively, which are significantly lower than those determined for the corresponding non-passivated monoliths. The SEM micrographs (Figure 7) of the cross sections of the passivated samples after oxidation at 1500 C can further prove their highly improved oxidation resistance as compared to the non-passivated samples. No cracks can be observed on the surface and within cross section of the passivated monoliths after oxidation. As shown in Figure 7, a dense, continuous oxide scale can be observed on the surface of each sample. Additionally, the isothermal oxidation behavior of the p(Hf0 .2Ta0 .8)C/SiC, p-(Hf0 .7Ta0.3)C/SiC, p-(Hf0 .9Ti0.1)C/SiC, and p-(Hf0.5Ti0.5)C/SiC was investigated at temperatures in the range from 1200 to 1500 C using an exposure time of 20 h. The oxidation curves are shown in Figure 8. Unlike the corresponding non-passivated monoliths, there is no detectable mass loss caused by the oxidation of segregated carbon and the mass gain after 20 h dwell at all temperatures is relatively low. The photographs of the oxidized samples (see Supporting Information, Figure S4) show that all the passivated monoliths are intact, and no cracks can be observed, which further proves the significant improvement of their oxidation resistance.  \\x0e  Figure 6. Oxidation curves of  the pre-oxidized monoliths during TGA assisted oxidation test  from 20 to 1500  C. Plot  in (b)  represents the  \\x0e  magnification of  the  plot  in (a);  isothermal oxidation part at 1500  \\x0e  grayish  area  in (b)  represents  the  C.  up to 50 mol%, the phase compositions reveals a mixture of HfO2 and HfTiO4; while for the TiO2 molar fractions \\x15 50 mol%, TiO2 and HfTiO4 co-exist.[72] The presence of ca. 50 mol% TiC solid-solution carbide phase is a reasonable explanation for the formation of HfTiO4 upon oxidation of p-(Hf0.5Ti0.5)C/SiC and also a suitable reason why there was no TiO2 present in the investigated samples. In conclusion, the formation of HfTiO4 is considered to be beneficial for the oxidation behavior of (Hf,Ti)  Figure 7. SEM images of  the fractured cross section of  the passivated monoliths after TG-oxidation test  from 20 to 1500  \\x0e  C.  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (8 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  Figure 8. Oxidation curves of  the passivated monoliths under isothermal oxidation test at 1200-1500  \\x0e  C for 20 h.  \\x0e  The values of the parabolic rate constants (Kp) of the passivated monoliths upon oxidation at temperatures from 1200 to 1500 C (see Table 3 \\x005-10 and Supporting \\x003 mg2/(cm4 h), and thus Information, Figure S5) are in the order of 10 are comparable to those measured for CVD SiC and sintered SiC, as mentioned above. Unlike the non-passivated samples having Kp values with a negative temperature evolution (i.e., the Kp (Figure 3 and ref. values of the passivated samples exhibit a positive temperature evolution (i.e., Kp1200 < Kp1300 < Kp1400 < Kp1500; Table 4). The values of the activation energy (Ea) for the oxidation processes of p-(Hf0.2Ta0.8)C/SiC, p-(Hf0.7Ta0.3)C/SiC, p-(Hf0.9Ti0.1)C/SiC, and  Kp1200>Kp1300>Kp1400>Kp1500)  [52]),  \\x0e  C are 199.0, 228.4,  p-(Hf0.5Ti0.5)C/SiC monoliths at 1200-1500 \\x001, respectively. 158.7, and 258.0 kJ mol Thus, the passivation of the investigated monolithic samples prior oxidation leads to a significant decrease of their apparent Kp values (3-4 orders of magnitude lower than those measured or the non-passivated samples). For instance, the Kp value of p-(Hf0.9Ti0.1) C (6.9 \\x01 10 \\x004 mg2/(cm4 h)) is the highest among the C/SiC at 1200 four investigated passivated monoliths; this correlates to the large molar fraction of HfC in the (Hf,Ti)C solid solution. However, (Hf0.9Ti0.1)C/SiC (Kp ¼ 0.85 mg2/(cm4 h), Figure S5], the Kp value of compared with the corresponding non-passivated sample, that is, p-(Hf0.9Ti0.1)C/SiC is 3 orders of magnitude lower.  \\x0e  Table 4. Apparent parabolic oxidation rate Kp of  the passivated monolithic samples (calculated cf. Equation (4)).  Sample  p-(Hf0.2Ta0.8)C/SiC  p-(Hf0.7Ta0.3)C/SiC  p-(Hf0.9Ti0.1)C/SiC  p-(Hf0.5Ti0.5)C/SiC  Apparent parabolic oxidation rate Kp [10  \\x005 mg2/(cm4 h)]  1300  \\x0e  C  18.70  9.25  90.8  10.3  1400  \\x0e  C  55.10  28.10  338  33.1  1500  \\x0e  C  67.60  50.40  515  99.6  1200  \\x0e  C  4.76  2.25  69.0  2.83  Activation energy Ea [kJ mol  \\x001]  199.0  228.4  158.7  258.0  Adv. Eng. Mater. 2019, 21, 1800879  1800879 (9 of 11)  © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim  \\x0c', 'www.advancedsciencenews.com  www.aem-journal.com  \\x0e  1400 C. Moreover, SEM images (see Figure S8, S9 in Supporting Information) prove that the monoliths are well protected by a dense, crack-free oxide scale.  4. Conclusions  In the present work, a synergistic strategy to improve the hightemperature oxidation behavior of HfC/SiC is presented. The strategy consists of a combination of alloying TaC/TiC to HfC and of applying a passivation step to the ceramic nanocomposites prior to their exposure to high-temperature oxidation conditions. It is shown that the passivation treatment activates the formation of silica, which is sluggish at moderate temperatures; whereas the alloying of HfC with TaC and TiC induces a significant improvement of the oxidation resistance via the formation of oxidic phases such as Hf6Ta2O17 and HfSiO4/HfTiO4, respectively, which are very beneficial for the high-temperature oxidation behavior of the studied ceramic nanocomposites.  Supporting Information  Supporting Information is available from Wiley Online Library or from the author.  Acknowledgements  Financial support from the German Science Foundation (DFG; Bonn, Germany) and from the R&D Convergence Program of MSIP (Ministry of Science, ICT and Future Planning) and NST (National Research Council of Science & Technology) of Republic of Korea (Grant: CMIP-13-4-KIMS) were gratefully acknowledged. The authors thanked Prof. Olivier Guillon (Forschungszentrum Jülich, Germany) and Dr. Koji Morita (NIMS, Tsukuba, Japan) for support with the spark plasma sintering of the samples. Claudia Fasel (TU Darmstadt, Germany) was acknowledged for the thermogravimetric measurements, Dr. Cong Zhou and Dr. Yan Lu (TU Darmstadt, Germany) for elemental analyses.  the passivated  Figure 9. Oxidation curves a) and parabolic plots b) of  monoliths after oxidation at 1400  C for 100 h.  \\x0e  The photographs (Figure S4, Supporting Information) and SEM images (Figure S6, S7, Supporting Information) of the passivated monoliths after oxidation at 1200 C for 20 h further substantiate the beneficial effect of passivation prior to oxidation process, revealing that the passivated monoliths after oxidation are still black and intact, and no cracks can be observed (unlike the non-passivated monoliths[52]).  \\x0e  Conflict of  Interest  The authors declare no conflict of  interest.  Keywords  hafnium carbide, high-temperature oxidation, passivation, solid solution, ultrahigh-temperature ceramic nanocomposites  3.5. Long-Term Isothermal Oxidation Test of  Monoliths at 1400  C  \\x0e  the Passivated  Final Version: September 19, 2018  Received: August 16, 2018  Published online: October 21, 2018  \\x0e  The passivated monoliths were exemplarily studied concerning their long-term oxidation behavior at 1400 C for 100 h (Figure 9). All studied samples show parabolic behavior and the apparent parabolic rate constants (Kp) of p-TaC/SiC, p-(Hf0.2Ta0.8)C/SiC, p(Hf0.7Ta0.3)C/SiC, p-(Hf0.9Ti0.1)C/SiC, and p-(Hf0.5Ti0.5)C/SiC 2.79 \\x02 10 \\x005, 7.97 \\x02 10 \\x005, 1.41 \\x02 10 \\x004, 4.91 \\x02 10 \\x004, 8.05 \\x02 10 were \\x005 mg2/(cm4 h), and respectively. Thus, all passivated monoliths exhibit excellent long-time oxidation resistance at  [1] H. O. Pierson, Handbook of Refractory Carbides & Nitrides: Properties,  Characteristics, Processing and Applications, Noyes Publications, USA  1996.  [2]  L. Feng, J. M. Kim, S. H. Lee, S. J. Park, J. Am. Ceram. Soc. 2016, 99,  1129.  [3] E.  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},{
  "_id": 245,
  "PDF": "Spark plasma sintering and characterization of ZrC-TiB 2 composites.pdf",
  "Text": "[\"Contents lists available at ScienceDirect  Ceramics International  jou rna l homepage : www . e lsev ie r .com / loca te /ce ram in t  Spark plasma sintering and characterization of ZrC-TiB2 composites  Ozden Ormanci Ozturk, Gultekin Goller  ⁎  Istanbul Technical University, Metallurgical and Materials Engineering Department, Maslak, 34469 Istanbul, Turkey  A R T I C L E  I N F O  Keywords:  Zirconium carbide  Titanium diboride  Spark plasma sintering  A B S T R A C T  ZrC-based composites were consolidated from ZrC and TiB2 powders by the Spark Plasma Sintering (SPS) technique at 1685 °C and 1700 °C for 300 s under 40-50-60 MPa. Densification, crystalline phases, micro structure, mechanical properties and oxidation behavior of  the composites were investigated. The sintered  bodies were  composed of  a  (Zr,Ti)C solid solution and a ZrB phase. The densification behaviors of  the  composites were improved by increasing the TiB2 content and applied pressure. The highest value of hardness, 21.64 GPa, was attained with the addition of 30 vol% TiB2. Oxidation tests were performed at 900 °C for 2 h and the formation of ZrO2, TiO2 and B2O3 phases were identified by using XRD.  1.  Introduction  Zirconium carbide  (ZrC)  is a potential  candidate  for ultra high temperature materials, such as cutting tools,  jet engine parts,  leading  edges of re-entry space aircraft due to its high melting point (3420 °C), (6.59 g/cm3), high hardness  low density  (25.5 GPa), high electrical (350-440 GPa) [1-6].  conductivity  and  high modulus  of  elasticity  However,  low sinterability due  to strong covalent bonding and low  self-diffusion rates restrict  the potential practical applications of  this  material  [2,7,8].  To  promote  densification,  carbides  are  generally  sintered with metalic  additives  such  as Ni, Mo,  and  their  alloys  [9,10]. However, metalic  additives used for  improving  sinterability  degrade the high temperature properties and corrosion resistance of  the material  [11]. To improve the densification of UHTCs the Spark  Plasma  Sintering  (SPS)  technique  is  a  promising  approach.  This  technique  employs  a  pulsed  direct  current which  passes  through  graphite punch rods and dies with a uniaxial pressure. Furthermore,  enhanced  densification,  refined microstructures  and  clean  grain  boundaries, which result in an overall improvement in the material's performance, have been reported by other working groups [12-14].  Low oxidation resistance is another limitation for high-temperature  applications of ZrC-based composites. When ZrC is oxidized to non protective ZrO2, approximately 650 °C [15].  the maximum workable temperature is  reported as  Incorporation of one or more  carbide/  boride phase  into the ZrC matrix  is  the most  common method of  enhancing its oxidation resistance and mechanical properties. SiC is  the  additive which is mainly used to  improve oxidation resistance  through the formation of a protective SiO2 layer at high temperatures, however there has been limited investigation on the influence of TiB2  additions on the properties of ZrC [12,16,17].  In this  study, ZrC-TiB2 produced by using the  composites with different  compositions  were  Spark Plasma  Sintering  technique  at  different  temperatures  and  pressures. As  a  result,  the  unexplored  effects of TiB2 additions on densification, microstructure, mechanical properties and oxidation resistance of ZrC are reported.  2. Experimental procedure  ZrC (H.C. Starck Corp. Grade B, AB134580), and TiB2 (H.C. Starck, Grade D) powders were used as starting materials. SEM observations  shown in Fig. 1(a) and (b)  reveal  the presence of grain size hetero geneities. The crystals have an angular shape and most of them have an  average grain size ranging from 1 to 5 µm.  The powders were weighed and ball-milled for 24 h in ethanol using  Si3N4 balls in polyethylene bottles and dried in a drying-oven at 100 °C. Throughout the text, the composites were labelled ZTi10, ZTi20 and  ZTi30 according to their increasing TiB2 content. For instance, ZTi10 represents the composites which contain 10 vol% TiB2. The mixed powders were loaded into a graphite die with an inner  diameter of 50 mm and consolidated with an SPS apparatus (SPS-7.40  MK-VII, SPS Syntex Inc.) at temperatures of 1685 °C and 1700 °C for  300 s with a heating rate of 2.5 °C/s under vacuum. Uniaxial pressures  of  40-50-60 MPa were  applied  respectively. During  the  sintering  process  a  graphite blanket was placed around the die  in order  to  minimize  the heat  loss. The  temperatures were measured with an  optical pyrometer  (Chino,  IR-AH)  focused on the  graphite die  and  linear  shrinkages of  the specimens during sintering were monitored  through the displacement of a punch rod.  http://dx.doi.org/10.1016/j.ceramint.2017.03.199  Received 29 January 2017; Received in revised form 28 March 2017; Accepted 29 March 2017  ⁎ Corresponding author.  E-mail address: goller@itu.edu.tr (G. Goller).  Ceramics International 43 (2017) 8475-8481  Available online 01 April 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  MARK  \\x0c\", 'Bulk  densities were  determined  by  the  Archimedes’ method.  Crystalline phases were identified by X-ray diffraction (XRD; MiniFlex, Rigaku Corp.) using Cu-Kα radiation. The microstructure  of the starting powders and polished surfaces of the sintered specimens  were  observed  with  scanning  electron  microscopy  (FESEM;  JSM7000F, JEOL Ltd.). Vickers hardness (HV) values were measured  under load a of 9.8 N.  In the oxidation studies, all  samples were placed in an alumina  crucible with a minimum contact area, so that the maximum area of the  samples  could be  exposed to oxidation. A MoSi2 furnace (Nabertherm C42) was used to heat the samples. Composites  resistance heated  were exposed to stagnant air at 900 °C for 2 h. Sample weights were  measured before and after oxidation and dimensions were measured in  order to calculate the surface area of test specimens.  3. Results and discussion  3.1. Crystalline phases  FactSage thermochemical software was used to simulate the reac tions during the sintering process. Taking into account our working  parameters, the reaction of ZrC with TiB2 produces of a ZrC-TiC solid solution and ZrB2 phases (Fig. 2). TiC and ZrC have the same cubic structure and can form a complete solid solution [18]. However, owing  to  the  difference  in  the  ionic  radii  of  Ti  and  Zr  and  their  low  diffusivities, high sintering temperatures are required to obtain com plete  solid  solutions  using  conventional  sintering. However,  low  sintering temperatures can be applied to produce ZrC-TiC composites  in the case of SPS, as reported by other working groups [19,20].  Fig.  3  shows  the XRD patterns  of  the  starting  powders  and  composites sintered at 1700 °C for 300 s under 50 MPa. All composite  samples  contain  ZrC  peaks while  they  do  not  contain  TiB2, also absent  as  thermodynamically  expected (Fig. 4). TiC peaks  are  in  the XRD patterns of all compositions. Fig. 5 illustrates the profiles of (2θ~33.07°)  the strongest XRD peaks  for different amounts of TiB2 the diffraction peaks shift to higher 2θ  additions. It can be seen that  degrees as the content of Ti  increases. Because the ionic radius of Zr  (0.98 Å) is larger than Ti (0.67 Å), the introduction of Ti atoms into ZrC  decreases  the lattice parameter. The calculated lattice parameters of  XRD patterns  are 4.68 Å, 4.67 Å and 4.66 Å for ZTi10, ZTi20 and  ZTi30 composites, respectively. Thus,  the shift of  (Zr,Ti)C diffraction  peaks towards larger angles can be observed. By the determination of  the lattice parameters of composites,  it is proved that substituted solid  solutions were formed.  Although the Factsage results indicate that the reaction of ZrC with  TiB2 generates a ZrB2 phase, a zirconium monoboride (ZrB) phase was observed instead of ZrB2 in the XRD patterns. According to the Zr-B phase diagram [21], a ZrB phase can also exist between the temperatures 1073-1523 K in addition to the most stable compounds of ZrB2. Although Portnoi and Romashov [22] failed to observe a ZrB phase in  Fig. 1. SEM images of starting powders of ZrC (a), and TiB2 (b).  Fig. 2. Factsage results of the reaction of ZrC with TiB2.  O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  8476  \\x0c', 'O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  Fig. 5. The profiles of the strongest XRD peaks in the ZrC phases for different amounts  of TiB2 additions.  Fig. 3. XRD patterns of starting powders and spark plasma sintered ZTi10 and ZTi30  composites.  Fig. 4. The equilibrium products of the reaction of ZrC with TiB2.  their studies, its existence was later confirmed by other working groups  [23]. Glaser and Post [24] found that the ZrB phase has an extremely  narrow range of stability and rapid quenching is necessary to stabilize it  at room temperature. However, other investigations have shown that  the ZrB phase is retained easily at room temperature on slow cooling.  The presence of impurities such as carbon, nitrogen and oxygen is also  considered as another stabilizing factor in the formation of ZrB due to  Fig. 6. Relationship between displacement,  temperature and the time dependence of  isothermal displacement of ZTi10 and ZTi20 composites  sintered at 1700 °C under  40 MPa (a), ZTi10 composites sintered at 1700 °C under 40-50-60 MPa (b).  and the presence of carbon impurities originating from the graphite  paper and die used.  the high reactivity of zirconium [23]. In this case, the existance of a ZrB  3.2. Densification behavior  phase can be attributed to the sintering temperature of the composites  The densification of  the  specimens during  the SPS process was  8477  \\x0c', 'O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  Fig. 7. SEM images of (a) ZTi10 sintered at 1685 °C under 50 MPa, (b) ZTi20 sintered at 1685 °C under 50 MPa, (c) ZTi10 sintered at 1685 °C under 60 MPa, (d) ZTi10 sintered at  1700 °C under 60 MPa.  Fig. 8. EDX mapping of ZTi20 composite sintered at 1685 °C, 60 MPa.  evaluated by the displacement of punch rods due to the shrinkage of  tively. The shrinkage started at a lower temperature (1500 °C) with the  the  composites.  Fig.  shrinkage  curves  at  6(a) displays 1400-1700 °C  the  effects  of  composition  and  the  time  dependence  on  of  addition of 30 vol% TiB2. Effect of solid solution formation on the densification of ZrC-TiC  isothermal displacements at 1700 °C for up to 300 s. The shrinkage  ceramics has been studied by Niu et. al. and it was reported that TiC  of ZTi10 started at 1560 °C and stopped at 1700 °C. The addition of  doping is extremely effective in promoting the densification of ZrC [25].  20 vol% TiB2 decreased the initiation and completion temperatures of shrinkage from 1560 °C to 1510 °C and 1700 °C to 1675 °C respec In  this  study,  the  improvment  of  densification  behavior  could  be  attributed to the  formation of  (Zr,Ti)C solid solution by  increasing  8478  \\x0c', 'TiB2 addition. The effect  of  sintering  pressure  on  densification  is  shown  in  Fig. 6(b). The shrinkage of  the ZTi20 composite sintered at 1700 °C  under  40 MPa  started  at  1550 °C  and  continued  up  to  1700 °C,  whereas  the  composite  sintered  at  60 MPa  started  to  shrink  at  1490 °C and stopped at 1670 °C.  Increment  in pressure  shows  the  same effect on ZTi10 and ZTi30 composites. It is reported that applied  pressure  increases  the number of  contact points between particles,  thereby  reduces  the  effective diffusion distance. Also,  increment  in  pressure leads to particle rearrangement, more effective packing with  the existing large agglomerates and destruction of agglomerates, which  yields better homogeneity and higher densities [26].  3.3. Microstructural observation  Microstructures of the polished surfaces of composites are shown in  Fig. 7. The distribution of grains is homogeneous and small entrapped  pores are found. Elemental mapping of the ZTi20 composite with EDX  confirmed a uniform distribution and showed that  the dark  areas  represent Ti-rich regions  and the  lighter  ones Zr-rich matrices  as  shown in Fig. 8.  Microstructures  of ZTi10 and ZTi20 sintered at 1685 °C under  50 MPa are shown in Fig. 7(a) and (b). The sample containing 10 vol%  TiB2 has smaller grain size than samples containing 20 vol% TiB2. A slight increase in temperature resulted in a significant change in the  microstructure as shown in Fig. 7(c) and (d). Detailed analyses of  the  micrographs of ZTi10 composites indicate that  the average grain size  increases aproximately from 3 µm to 7 µm as the sintering temperature  increases from 1685 °C to 1700 °C respectively. In addition, the effect  of pressure on average grain size can be seen in Fig. 7(a) and (c). It  should be noted that  the applied pressure had no influence on grain  size and this finding is compatible with others in the literature [13,27].  3.4. Oxidation behavior  The oxidation behaviors of  the composites mainly depend on the  oxidation behaviors of the individual components. ZrC starts to oxidize  into the non protective and porous ZrO2 phase (Eq. (1)) at temperature above 600 °C and two processes are reported: a parabolic diffusion  reaction involving the partial replacement of  interstitial carbon in the  ZrC lattice with oxygen, and a linear surface reaction occuring at  the  ZrC-ZrO2 phase boundary [15].  ZrC (s)+2O2 (g)=ZrO2 (s)+CO2 (g)  (1)  Meerson et al. [28] studied the oxidation of ZrB2 and suggested the presence of ZrO covered by B2O3, during the oxidation reaction B2O3 evaporates and the oxygen reacts with ZrO to form ZrO2 (Eq. (2)).  ZrB2(s)+5/2 O2=ZrO2 (s)+B2O3 (l)  (2)  Shimada studied the oxidation behavior of commercial TiC powder  under non-isothermal conditions and reported the stages of oxidation.  According to the results, TiO2 was with increasing temperature (Eq. (3))  formed and increased in content  [29]. The grains cracked along  the  edges,  and the  cracks became  greater with the progression of  oxidation,  some of  the  grains were  completely broken into several  pieces.  TiC (s)+2O2 (g)=TiO2 (s)+CO2 (g)  (3)  Fig. 9 shows the XRD diffraction pattern of  the ZTi20 composite  after oxidation testing at 900 °C for 2 h and characteristic peaks of  ZrO2, TiO2 and B2O3 were identified. Pictures of the samples after oxidation testing are given in Fig. 10  for comparison. A slight increase in mass was observed for all samples  due  to the  formation of oxidation products. The mass  changes per  surface area of the ZTi10, ZTi20, ZTi30 composites sintered at 1700 °C  under  50 MPa  are  0.0964,  0.0712  and  0.0307  respectively.  The  composite prepared by adding 30 vol% TiB2 has the lowest weight gain. Fig. 11 shows the SEM images of the surface morphology of the  samples oxidized at 900 °C. The dark and light areas represent Ti-rich  and Zr-rich regions,  respectively. The big cracks and spallations of  ZTi10 composite can be seen in Fig. 11(a). The occurrence of the cracks  may be based on increase in volume due to the formation of oxidation  products  at pores. Less  cracks  are  formed in the ZTi20 composite  (Fig. 11(b)) and cracks could barely be observed in the composite of  ZTi30 (Fig. 11(c)). The  superior  oxidation resistance  of  the ZTi30  composite may be due to the increasing boride content.  3.5. Mechanical properties  The Vickers hardness values of ZrC-based composites at a load of  9.8 N are shown in Table 1, and Fig. 12 demonstrates the effects of TiB2 additions and sintering pressure on the hardness values of composites.  The  hardnesses  of  the  composites  increased with  increasing TiB2  Fig. 9. XRD pattern of spark plasma sintered ZTi20 composite after oxidation test at  900 °C for 2 h.  Fig. 10. View of the samples after oxidation test at 900 °C for 2 h.  O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  8479  \\x0c', 'addition and ZTi30  composite  showed  higher  hardness  than both  ZTi10  and  ZTi20  composites.  The  hardness  increased  from  17.02 GPa to 21.64 GPa when the TiB2 content 30 vol%. The increase in hardness appeared to be the result of  increased from 10 to  the  higher  theoretical  hardnesses  of  sintering  products  and  enhanced  densification with the formation of (Zr,Ti)C solid solution [30].  4. Conclusions  ZrC-based  composites  were  consolidated  from ZrC  and  TiB2 reaction between ZrC and TiB2 solution, which improved densification  powders by  the SPS method. The  produced  a  ZrC-TiC  solid  probably due to the volumetric diffusion of Ti and Zr atoms, and a  ZrB phase.  Increase  in pressure also enhanced the densification by  means of  increasing the number of contact points between particles  and reducing the effective diffusion distance. The hardness values of  the composites  increased with increasing TiB2 value of hardness, 21.64 GPa, was achieved in the ZTi30 composite  content. The highest  Fig. 11. SEM images of the surface morphology of oxidized ZTi10 (a), ZTi20 (b) and ZTi30 (c) samples.  Table 1  Compositions, sintering consitions, densities and hardness of spark plasma sintered specimens.  Code  Starting materials  Composite  Sintering temperature (°C)  Applied pressure (MPa)  Density  Hardness  ZrC (vol%)  TiB2 (vol%)  (g/cm3)  (GPa)  ZTi10  90  10  ZrC-TiC-ZrB2  1685  40  6.32  17.04 ± 0.34  50  6.36  17.93 ± 0.35  60  6.38  18.07 ± 0.32  1700  40  6.39  18.09 ± 0.21  50  6.42  18.11 ± 0.16  60  6.44  17.02 ± 0.35  ZTi20  80  20  ZrC-TiC-ZrB2  1685  40  6.11  19.72 ± 0.76  50  6.16  19.64 ± 0.44  60  6.14  19.68 ± 0.58  1700  40  6.17  19.59 ± 0.62  50  6.14  19.55 ± 0.58  60  6.18  19.76 ± 0.80  ZTi30  70  30  ZrC-TiC-ZrB2  1685  40  5.98  20.84 ± 0.36  50  6.01  20,75 ± 0.38  60  6.04  20,92 ± 0.45  1700  40  6.05  21.43 ± 0.56  50  6.09  21.56 ± 0.36  60  6.10  21.64 ± 0.27  Fig. 12. Effect of TiB2 additions and pressure on the hardness of ZrC-based composites.  O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  8480  \\x0c', 'O. Ormanci Ozturk, G. Goller  Ceramics International 43 (2017) 8475-8481  [12] V. Medri, F. Monteverde, A. Balbo, A. Bellosi, Comparison of ZrB2-ZrC-SiC composites fabricated by spark plasma sintering and hot-pressing, Adv. Eng. Mater. 7 (2005) 159-163. Z.A. Munir, U. Anselmi-Tamburini, The effect of electric field and pressure on the  [13]  synthesis and consolidation of materials: a review of the spark plasma sintering method, J. Mater. Sci. 41 (2006) 763-777. [14] D.S. Perera, M. Tokita, S. Moricca, Comparative study of  fabrication of Si3N4/SiC composites by spark plasma sintering and hot isostatic pressing, J. Eur. Ceram. Soc. 18 (1998) 401-404. [15] A.K. Kuriakose, J.L. Margrave, The oxidation kinetics of zirconium diboride and zirconium carbide at high temperatures, J. Electrochem. Soc. 111 (1964) 827-831. I. Akin, G. Goller, Mechanical and oxidation behavior of spark plasma sintered ZrB2-ZrC-SiC composites, J. Ceram. Soc. Jpn. 120 (2012) 143-149. [17] D. Pizon, L. Charpentier, R. Lucasa, S. Foucaud, A. Maître, M. Balat-Pichelin, Oxidation behavior of spark plasma sintered ZrC-SiC composites obtained from the polymer-derived ceramics route, Ceram. Int. 40 (2014) 5025-5031. [18] R.F. Bunshah, R. Nimmagadda, W. Dunford, Structure and properties of refractory  [16]  [20]  compounds deposited by electron beam evaporation, Thin Solid Films 54 (1978) 85-106. [19] R.B. Acicbe, G. Goller, Densification behavior and mechanical properties of spark plasma-sintered ZrC-TiC and ZrC-TiC-CNT composites, J. Mater. Sci. 48 (2013) 2388-2393. Y. Li, H. Katsui, T. Goto, Spark plasma sintering of ZrC-TiC composites, Ceram. Int. 41 (2015) 7103-7108. [21] K.I. Portnoi, V.M. Romashov, L.N. Burobina, Constitution diagram of the system zirconium - boron, Sov. Powder Metall.+ 9 (1970) 577-580. [22] K.I. Portnoi, V.M. Romashov, Binary constitution diagrams of systems composed of various elements and boron - a review, Sov. Powder Metall.+ 11 (1972) 48-56. Y. Champion, S. Hagege, Structural analysis of phases and heterophase interfaces in the zirconium-boron system, J. Mater. Sci. 33 (1998) 4035-4041. F.W. Glaser, B. Post, Phase diagram zirconium-boron, Trans. Metall. Soc. AIME 197 (1953) 1117-1118. [25] B. Niu, F. Zhang, W. Ji, J.Y. Zhang, Z.Y. Fu, W.M. Wang, Effect of solid solution  [23]  [24]  formation on densification of spark plasma sintered ZrC ceramics with tic as sintering aid, Adv. Appl. Ceram. 115 (2015) 55-59. P. Barick, D. Chakravarty, B.P. Saha, R. Mitra, S.V. Joshi, Effect of pressure and  [26]  temperature on densification, microstructure and mechanical properties of spark plasma sintered silicon carbide processed with β-silicon carbide nano powder and sintering additives, Ceram. Int. 42 (2016) 3836-3848. [27] D.V. Quach, H. Avila-Paredes, S. Kim, M. Martin, Z.A. Munir, Pressure effects and  grain growth kinetics in the consolidation of nanostructured fully stabilized zirconia by pulsed electric current sintering, Acta Mater. 58 (2010) 5022-5030. [28] G.A. Meerson, A.F. Gorbunov, Activated sintering of zirconium boride, Inorg. Mater. 4 (1968) 267-270. S. Shimada, A thermoanalytical study of oxidation of TiC by simultaneous TGADTA-MS analysis, J. Mater. Sci. 31 (1996) 673-677. [30] D.L. Yung, S. Cygan, M. Antonov, L. Jaworska, I. Hussainova, Ultra high-pressure  [29]  spark plasma sintered ZrC-Mo and ZrC-TiC composites, Int. J. Refract. Met. Hard Mater. 61 (2016) 201-206.  which was prepared with a 30 vol% TiB2 addition. The ZTi30 composite also exhibited the highest oxidation resistance due to the increasing  boride content.  Acknowledgments  This work was supported by the Scientific Research Project Funds  of Istanbul Technical University (Project Number: 38388). The authors  thank H.H. Sezer for SEM studies. The authors would also like to thank  to Mimar Sinan Fine Arts University, Material Research Center  for  Cultural Property and Artworks for the infrastructure support.  References  [3]  [1]  S.Q. Guo, Y. Kagawa, T. Nishimura, D. Chung, J.M. Yang, Mechanical and physical behavior of spark plasma sintered ZrC-ZrB2-SiC composites, J. Eur. Ceram. Soc. 28 (2008) 1279-1285. [2] D. Sciti, S. Guicciardi, M. Nygren, Spark plasma sintering and mechanical behaviour of ZrC-based composites, Scr. Mater. 59 (2008) 638-641. S. Sagdic, G. Goller, Densification behavior and mechanical properties of spark plasma sintered ZrC-SiC and ZrC-SiC-CNT composites, J. Aust. Ceram. Soc. 50 (2014) 76-82. J. Sha, J. Li, S. Wang, Z. Zhang, Y. Wang, J. Dai, Microstructure and mechanical properties of hot-pressed ZrC-Ti-CNTs composites, Mater. Des. 107 (2016) 520-528. J. Li, Z. Zhang, S. Wang, Y. Wang, J. Dai, Y. Zu, J. Sha, Densification and  [4]  [5]  characterization of hot-pressed ZrC-based composite doped with Nb and CNT, Mater. Des. 104 (2016) 43-50. [6] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, Properties,  Characteristics, Processing and Applications, Noyes Publications, New Jersey,  1996.  [7] B. Ma, X. Zhang, J. Han, W. Han, Fabrication of hot-pressed ZrC-based composites, J. Aerosp. Eng. 223 (2009) 1153-1157. L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Pressureless sintering of ZrC-based  [8]  ceramics by enhancing powder sinterability, Int. J. Refract. Met. Hard Mater. 29 (2011) 516-521. T. Suzuki, H. Matsumoto, N. Nomura, S. Hanada, Microstructures and fracture  [9]  [10]  toughness of directionally solidified Mo-ZrC eutectic composites, Sci. Tech. Adv. Mater. 3 (2002) 137-143. J.Y. Ko, S.Y. Park, D.Y. Yoon, S.J.L. Kang, Migration of intergranular liquid films and formation of core-shell grains in sintered TiC-Ni bonded to WC-Ni, J. Am. Ceram. Soc. 87 (2004) 2262-2267. [11] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. Zhou, Ultra-High Temperature  Ceramics, Materials for Extreme Environment Applications, John Wiley, New  Jersey.  8481  \\x0c']"
},{
  "_id": 246,
  "PDF": "Stability of ultra-high-temperature ZrB2–SiC ceramics under simulated atmospheric re-entry conditions.pdf",
  "Text": "['Journal of the European Ceramic Society 27 (2007) 4797-4805  Stability of ultra-high-temperature ZrB2-SiC ceramics under simulated atmospheric re-entry conditions  Fr ´ed ´eric Monteverde a,∗  , Raffaele Savino b  a ISTEC, Institute of Science and Technology for Ceramics, National Research Council, Via Granarolo 64, 48018 Faenza, Italy b DIAS, Dipartimento di Ingegneria Aerospaziale, University of Naples “Federico II”, P.le V. Tecchio 80, 80125 Naples, Italy  Received 12 October 2006; received in revised form 13 February 2007; accepted 18 February 2007  Available online 29 March 2007  Abstract  Microstructure modifications of an ultra-high temperature ZrB2 -SiC ceramic exposed to ground simulated atmospheric re-entry conditions were investigated and discussed. Fluid dynamic numerical calculations were carried out to correlate and explain the experimental results. The crosssectioning of the ceramic models after exposure (examined by SEM) showed a compact scale of zirconia (20 \\u242em thick) underlying an external silica thin coating. A partially SiC-depleted region, a few microns thick, underneath the zirconia sub-scale was also seen. The post-test analyses confirmed the potential of the ZrB2 -SiC composite to endure re-entry conditions with temperature approaching 2000 C, thanks to the formation of a steady-state external multiphase oxide scale. Numerical calculations, which simulated the chemical non-equilibrium flow around the ceramic model, matched well the experimental results only assuming a very low catalytic surface behavior. © 2007 Elsevier Ltd. All rights reserved.     Keywords: Structural applications; Thermal properties; ZrB2 ; SiC  1.  Introduction  Improved interest in ultra-high temperature ceramics (UHTCs) is being animating the scientific community over the past decade.1-4 This emerging attention is driven by the demand of developing re-usable hot structures as thermal protection systems (TPS) of space vehicles able to re-enter from Low Earth Orbit (LEO) at relatively high speed (order of 8 km/s). In contrast to traditional blunt capsules or Shuttle-type vehicles, characterised by poor gliding capabilities and complex TPS, the future use of UHTCs opens new prospects for the development of space planes with slender fuselage noses and sharp wing leading edges. In fact, sharp geometries imply peak heat fluxes in the order of 1 MW/m2 that state-of-art hot structures such as SiC-coated C-C composites are not able to withstand.1,4,5 The use of UHTCs would also imply lower aerodynamic drag, improved flight performances and crew safety, due to the larger cross range and manoeuvrability along more gentle re-entry trajectories.3,6,7  ∗  Corresponding author. Tel.: +39 0546 699758; fax: +39 0546 46381.  E-mail address: fmonte@istec.cnr.it (F. Monteverde).  0955-2219/$ - see front matter © 2007 Elsevier Ltd. All rights reserved.  doi:10.1016/j.jeurceramsoc.2007.02.201     IVA group transition metal diborides such as ZrB2 and HfB2 have been indicated as promising candidate materials for use in these aerospace applications, primarily for melting temperatures greater than 3200 C. Other favorable characteristics include high elastic modulus, high thermal conductivity, retained strength at elevated temperature, relatively good thermal shock resistance and modest thermal expansion.8 However, pure metal diborides do not fully possess oxidation resistance necessary for surviving reliably the oxidising environment typical of hypersonic re-entry into the Earth atmosphere. Expose HfB2 to air at elevated temperatures, for instance, leads to form HfO2 and liquid B2O3 : above 1600 C, due to the disruption of the HfO2 scale when B2O3 starts boiling very actively,9 HfB2 substantially deteriorates. A special family of diborides-based composites have demonstrated to overcome this limitation, thanks to the incorporation of SiC which improves mechanical properties10,11 and resistance to oxidation.1,3,5,9 From a technological point of view, the effective densification of ZrB2 , due to strong covalent bondings and low selfdiffusion coefficients, typically requires sintering temperatures above 2000 C and applied pressures in atmosphere-controlled furnaces.12 Instead, several studies demonstrated that SiC added to ZrB2 enhanced sinter ability and mechanical properties,10,11        \\x0c', '4798  F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805           in addition to the resistance to oxidation1,13-15 of the diboride alone. For ZrB2 oxidised in air at elevated temperatures, ZrO2 and liquid B2O3 are formed.14 Below 1200 C, molten B2O3 glass is basically retained in the porous ZrO2 structure due to high wettability and considerable surface tension. However, appreciable volatilisation of B2O3 starts taking place above 1200 C,16 leading to the formation of scarcely protective porous ZrO2 scale. The addition of SiC to ZrB2 (for temperatures above 1200 C) significantly reduces the oxidation rates, thanks to the stability, at much higher temperatures than for the pure diboride, of a protective silica-based glassy layer covering the surfaces exposed to air. Passive-to-active transition of the SiC oxidation has been also observed, resulting in regions beneath the oxide scales formed upon the ZrB2-SiC base ceramic partially depleted of SiC.10,14,17 How UHTCs behave when exposed to convective flows with significant concentrations of atomic oxygen (as is the case of shock waves in hypersonic flows) is a central question that has not received much attention to date. Indeed, a very few number of (expensive) large-scale arc-jet facilities are available for testing UHTCs (and other materials) in hypersonic flows with specific total enthalpies in the order of 10-20 MJ/kg. Thus, the majority of the oxidation studies on UHTCs employed conventional airfurnaces, which ultimately fail to expose materials to significant levels of atomic oxygen. In this study, microstructure modifications of ZrB2 -SiC composite subjected to arc-jet testing were investigated and discussed. Fluid dynamic numerical calculations were carried out to correlate and explain the experimental results.  2. Experimental  2.1. Material processing          A ZrB2-SiC ceramic was prepared from commercially available powders of ZrB2 (grade B, H.C. Starck, Germany) and SiC (BF12, H.C. Starck, Germany). The powder mixture of ZrB2 + 15 vol.% SiC was ball-milled in ethanol for 1 day using hard milling media, dried with a rotating evaporator, and sieved through a mesh screen with 250 \\u242em openings to minimize agglomeration. The powder blend was then uniaxially hotpressed in an inductively excited BN-lined graphite die, heating up to 1820 C (about 20 C/min average heating rate and 15 min isothermal stage at 1820 C). The temperature was measured by means of an optical pyrometer focused on the graphite die. Bulk density and theoretical density were evaluated using the Archimedes’ method (water as immersing medium) and the rule-of-mixture, respectively. The relative density was calculated dividing the bulk density by the theoretical density. The phase composition analysis was carried out with a scanning electron microscope (SEM, mod. S360, Leica Cambridge, UK) combined with an energy dispersive X-ray microanalyser (EDS, mod. INCA Energy 300, Oxford Instruments, UK). Polished sections of the as-sintered material were prepared with successively finer diamond-based abrasives ranging from 50 to 0.25 \\u242em.  Fig. 1. Hemispheric ZrB2 -SiC model used for arc-jet testing, curvature radius R = 7.5 mm.  Thermal diffusivity (DTH ) was measured through the laser flash method (mod. LFA-427, NETZSCH Ger ¨atebau GmbH, Germany), while heat capacity (CP ) using a modulated differential scanning calorimetry (mod. MDSC, TA Instruments, USA). Thermal conductivity (KTH ) was estimated through the expression KTH = DTHCPρ, ρ is the bulk density.  2.2. Plasma torch testing  ZrB2 -SiC models with a hemispheric shape (Fig. 1) were machined through diamond-loaded tools, and then exposed to sustained enthalpy flows using the arc-jet facility available at the University of Naples. The facility is equipped with a 80 kW plasma torch that operates in inert gas (He, N2 , Ar and their mixtures) at mass flow rates up to 5 g/s. Infrared and optical windows in the test chamber allow visual inspection and diagnostic analyses. An automatic control system monitors the main parameters of the apparatus (voltage and current of the arc heater, water cooling temperature, mass flow rate). In particular, the specific total enthalpy (H) was evaluated through an energy balance between the energy supplied to the gas by the arc heater and the energy transferred to the cooling system (measured by the water temperature jump between inlet and outlet). The output data, processed via a dedicated software, allow not only the evaluation of the surface temperature profile versus exposure time of the model, but also some numerical-experimental correlations. Due to the extremely high thermal loading upon the ceramic models, surface chemical reactions like oxidation can be responsible for changes in the material’s emissivity. To overcome this problem, the experiments were carried out with a radiation ratio pyrometer (Infratherm ISQ5, Impac Electronic Gmbh, Germany) which operates both in two colour and in the single colour function. In the two colour mode the instrument makes use of the ratio of two spectral radiances, measured at different  \\x0c', 'F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  4799  Fig. 2. Fracture (a) and polished surface (b) of ZrB2 -SiC, SEM micrographs: round dark features are SiC (b).  wavelengths (0.9-1.05 \\u242em), to evaluate the true temperature. This overcomes the problem of the emissivity knowledge since it is supposed to be the same at both wavelengths. Once the temperature had been measured with the ratio pyrometer, its value was input to evaluate the spectral emissivity using the single colour function (λ = 0.9 \\u242em). In combination with the pyrometer, an infrared thermo-camera (Thermacam SC 3000, FLIR Systems, USA) was used to measure the surface temperature distributions and the spectral emissivity in the long wave range of the thermograph (λ = 9 \\u242em). The ceramic models exposed to the hot stream were further analysed by SEM-EDS on surface and after cross-sectioning. A thin carbon coating was deposited on the analysed surfaces to prevent charging effects.  3. Numerical simulation of the plasma torch flow  To assess if the environment generated by the plasma torch at atmospheric pressure conditions was able to reproduce heat fluxes, temperatures and reactive environments typical of atmospheric re-entry, a systematic numerical analysis has been carried out. The computations have been carried out solving the full Navier-Stokes equations for a turbulent multireacting gas mixture with five chemical species (O, O2 , NO, N, N2 ) in chemical non-equilibrium. Each species of the mixture was assumed to behave as a thermally perfect gas, where translational-rotational and vibrational-electronic degrees of freedom were characterised by different temperatures. Vibrational-translational energy exchanges have been modeled according to the Landau-Teller model, while the vibrational relaxation time was derived from the Millikan-White formula, with Park correction for high temperatures.18 Chemical and vibrational non-equilibrium is taken into account using the Park model.19-21 The fluid-dynamics equations have been numerically solved in the computational domain (plasma torch and test chamber). Convective fluxes were computed according to Roe’s Flux Difference Splitting scheme. Integration of the equations was performed implicit in time, until steady state was achieved, solving the linearised system of equation by the multigrid technique.  4. Results and discussion  4.1. Base microstructure  The bulk density of the sintered ZrB2-SiC pellets was 5.6 ± 0.02 g/cm3 , that corresponds to a relative density higher than 99%. The example of a fracture surface in Fig. 2a reveals the typical grain structure, with regular ZrB2 grains of about 2 \\u242em average size. The polished section examined by SEM (Fig. 2b) confirmed that the present composite is fully dense (i.e. no apparent porosity). In addition, the investigation by SEM provided evidence of SiC dispersed intergranularly within the ZrB2 matrix, sometimes in agglomerates (maximum size 5 \\u242em). Detailed SEM-EDS inspections of the microstructure allowed for highlighting also some secondary reaction products in very limited amounts like zirconia and a glass phase among aggregated SiC particles (Fig. 3). A more extensive description is reported elsewhere.22 The thermal conductivity (KTH ) varied from 62 to 65 W/m K over the temperature range 30-1200 C. Higher KTH values from 104 to 76.2 W/m K over the temperature range 30-1200 C were an UHTC composed of ZrB2 + 20 vol.% SiC.23 reported for        Fig. 3. Polished section of ZrB2 -SiC, SEM micrograph: glass phase (white arrow) and zirconia (black arrow) are marked.  \\x0c', '4800  F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  Fig. 4. Computed mass fractions of the different species along the torch axis  (x).  A variation in KTH was system.3  likewise observed in the HfB2-SiC  4.2. Plasma flow characterisation  The numerical simulations herein presented refer to a Nitrogen plasma jet with mass flow rate of 1 g/s, for a power transferred to the fluid of 15 kW: mass fraction profiles of different species along the symmetry axis of the torch (Fig. 4), and specific total enthalpy along the symmetry axis of the exhaust jet (Fig. 5) are shown. At the exit of the torch the partially dissociated Nitrogen expands through a nozzle (5 mm in diameter), comes into contact with the surrounding air at ambient conditions, so that oxygen in the atmosphere dissociates and a reacting mixture composed by O2 , N2 , NO, O, N is formed (Fig. 4). The specific total enthalpy decreases rapidly along the torch axis (Fig. 5). Changing the distance between the torch exit and the specimen, different conditions can be established. In  Fig. 5. Computed specific total enthalpy H and specific chemical enthalpy hd along the torch axis x, x = 0 the torch outlet.  Fig. 6. Stagnation point heat flux (q) computed at different distances from the  exit torch (x). The continuous lines correspond to the engineering formulas (EF)  for the different assumptions of fully catalytic (FC) and non catalytic (NC) wall.  The two points are evaluated with computation fluid dynamics (CFD).  particular, the range of interest is between 4 and 12 cm, with a corresponding decrease of the specific total enthalpy from about 18 to 3 MJ/kg. Fig. 5 also shows the contribution of the specific chemical enthalpy (hd ) to the specific total enthalpy due to the dissociated chemical species. The numerical results have been interpolated to show that the here reported semi-empirical formulas found for the stagnation point heat flux (q) at hypersonic conditions24  (cid:2)  (cid:2)  ·  q  FC = 2.75 × 10  −5  = 2.75 × 10  −5  (H0 − hw )1.17 ,  ·  q  NC  (H0 − hd − hw )1.17  (1)  p  R  p  R  are still valid in these conditions, where the NC and FC subscripts refer to a non catalytic wall and to a fully catalytic wall, respectively, H0 is the specific total enthalpy at the exit of the torch, hd is the chemical enthalpy stored in the atomic species; hw is the wall enthalpy, R the gas constant and p the stagnation point pressure. In Fig. 6 the stagnation point heat fluxes evaluated with the semi-empirical formulas are compared to the computational fluid dynamics (CFD) results and a good agreement is found.  4.3. Arc-jet testing  The present ZrB2 -SiC composite was subjected to arc-jet testing which, as verified in Section 4.2, enabled a ground simulation of conditions expected during an atmospheric re-entry. In particular, the hemispheric models (Fig. 1) were exposed to a reactive gas mixture with an average total enthalpy in the order of 10-20 MJ/kg. Fig. 7a shows the temperature profiles of the model’s surface during two arc-jet tests. During the test numbered 1 in Fig. 7a, the average specific total enthalpy was 17.3 MJ/kg during the first 60 s; the arc current was then increased and the specific total enthalpy reached a value of 19.7 MJ/kg and was maintained for approximately 270 s. In the test numbered 2 in Fig. 7a, a similar procedure was applied and the specific total enthalpy was 14.8 MJ/kg during the first 60 s and 16.6 MJ/kg for about 1 min. The corresponding values of  \\x0c', 'F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  4801  Fig. 7. Temperature (T) profiles vs. time during arc-jet testing of the ZrB2 -SiC model (a), and its post-test appearance (b); specific total enthalpies during test 1 (H11 and H12 ) and test 2 (H21 and H22 ) are indicated.  Fig. 8. Surface of the ZrB2 -SiC model after arc-jet testing (refer to the curve 1 in Fig. 7a), SEM micrographs.           the stagnation point heat fluxes, computed by numerical simulations, are in the order of 8-10 MW/m2 . In both cases, a value of about 0.9 was found for the emissivity at the highest temperatures. Just for comparison, arc-jet tests performed on ZrB2-SiC and HfB2 -SiC materials elsewhere described3,14 involved stagnation point heat fluxes in the order of 3.5-4 MW/m2 . With reference to the test numbered 1 in Fig. 7a, the surface temperature of the (hemispheric) model rose during exposure up to 1850 C within about 1 min. Increasing the specific total enthalpy up to about 20 MJ/kg, the surface temperature rapidly reached a value of about 1930 C, followed by a slow decay to 1910 C for the remaining duration of the test (i.e. 5 min). After exposure, the specimen appeared slightly glazed, but maintained an apparent integrity: a post-test picture of the model is displayed in Fig. 7b. Analyses by SEM-EDS of the exposed surface revealed a coherent and smooth silica glass coating decorated with partly aggregated zirconia crystals of various size and shape (Fig. 8). Silicon and oxygen apart, neither zirconium nor boron were detected in the (outermost part of the) external glass. To have a more detailed insight into the microstructure evolution upon the model’s surface during exposure, a specimen was tested at very similar conditions (specific total enthalpy 16.6 MJ/kg and total duration of about 2 min, see curve 2 in Fig. 7b). Though the visual appearance does not bring on significant differences compared to the test numbered 1 (Fig. 7a) early described, the model’s surface after this exposure revealed  clustered craters at the top of a glass coating (Fig. 9). The typical aggregations of zirconia crystals previously emphasised in Fig. 8 were not observed. The cross-sectioning and polishing of the models after arcjet testing provided insight into microstructure details of great interest. The elemental mapping by SEM-EDS throughout the oxidised cross-section revealed the formation of a multiphase layered structure (Figs. 10 and 11). The outermost layer, a few microns thick, consists of silica, and adheres to an oxide sub-scale, 20 \\u242em thick. The lastly cited scale includes zirconia  Fig. 9. Surface of the ZrB2 -SiC model after arc-jet testing (refer to the curve 2 in Fig. 7a), SEM micrograph.  \\x0c', '4802  F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  The SEM-EDS analysis of the cross-sectioned specimen associated to curve 2 in Fig. 7a (duration of about 2 min) presented very similar details. However, different from the ZrB2-SiC model arc-jet tested according to curve 1 in Fig. 7a, the occurrence of a pitted surface was linked to reactions involving the evolution of gaseous by-products. In fact, during the initial exposure to such harsh convectively heating conditions, SiO, B2O3 and other highly volatile boron sub-oxides, which tend to evolve outside, most likely coalesced in bubbles inside the external forming glass. Shear forces associated to the hot stream facilitated the bursting of bubbles and, for prolonged exposure (curve 1 in Fig. 7a), the smoothing of the outermost glassy coating. According to the SEM-EDS analytical resolution, silicon and oxygen apart, the characteristic peaks for boron and zirconium, due to the very low intensities, were of uncertain attribution. The postulation of mass loss mechanisms were not correlated to quantitative evaluations of weight changes and size recession of the models. The arc-jet tests presented in Fig. 7a provided an indication for the potential of this material for sharp leading edge applications: the significance of the microstructure changes after exposure to (simulated) re-entry conditions resulted within acceptable limits. It is the authors opinion that reasons for the configurational stability coupled to the high oxidation resistance of the present UHTC when subjected to heat fluxes typical of a re-entry environment should be inquired into the coherence of the passivating oxide scales covering its surfaces. The multiphase oxide scales of the present UHTC maintain protective capabilities even at extremely high temperatures thanks to the prompt recovery after bubble bursts: in that case, viscosity of the glass as well as its wetting behavior to the underlying oxide scale played a fundamental role. When the model’s temperature levelled off at about 1930 C (test 1 in Fig. 7a), the passivating silica layer which formed from the oxidation of SiC supported the maintenance of a steady state temperature. In addition, the condensation of volatile SiO (and boron sub-oxides) at the base of the oxide sub-scale (Fig. 10) could explain the stability of that microstructure as a quasi steady-state region, volatilising from the outer surface and re-constituting at lower depth. The crystalline zirconia that appears adherent to the formed glass (Figs. 10 and 11) suggests an extremely low solubility between the cited couple of compounds. Due to the preferential evaporation of boron oxides, precipitates of zirconia decorating the outermost glass layer after a relatively long exposure may indicate a change in ZrO2 solubility for borosilicate glasses versus the pure B2O3 . The improved performances of this material under strong heat flux conditions are deemed to arise also from its high thermal conductivity. This requisite allows heat to be conducted from the stagnation points to colder zones, and from there re-radiated away. In this respect, this class of UHTCs differs a typical insulating thermal protection structure which rejects heat almost solely by radiation. In addition, surface emissivities at very high temperatures around 0.9 helped the material maintain steady state surface temperature, levelling off for instance around 1930 C as in the test early described in Fig. 7a.        Fig. 10. Cross-section of the ZrB2 -SiC model after arc-jet testing (refer to the curve 1 in Fig. 7a), SEM micrograph; the oxide sub-scale (OX) underlying the  external glass layer (GL), and the SiC-depleted region (DE) close to the virgin  bulk (BU) are shown.  crystals embedded within a glass whose composition does not differ significantly from that located on the surface. Underneath, a layer partly depleted of SiC (10 \\u242em thick) has formed, and underlines the un-altered ZrB2-SiC matrix. The creation of such an inner porosity most likely was caused by the active oxidation of SiC, which represents a well-known phenomenon in SiC and SiC-containing materials.25-26 The interface between the fully oxidised scale and that partially depleted of SiC appears sharp, i.e. any transitional regions have formed. The external glass coverage apart, analogous multilayered structures were observed in MB2 -SiC models subjected to arc-jet testing.3,14,27  Fig. 11. Cross-section of the ZrB2 -SiC model after arc-jet testing (refer to the curve 1 in Fig. 7a, SEM micrograph by back-scattered electrons) and associated  elemental maps by SEM-EDS.  \\x0c', 'F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  4803     oxygen,10,13-15,17,22  temperature, found dramatically enhanced oxidation rates on Si, SiC and Si3N4 in dissociated versus molecular oxygen, under the same temperature and pressure.28 As the oxide scales development in a oxidising (re-entry) environment is concerned, it actually involves the composite interaction of factors primarily associated with mass transport. In contrast to oxidation studies in conventional high temperature furnaces where the primary oxidant is the molecular the transport mechanisms governing oxidation under an oxygen atom exposure are not yet clearly identified, even in pure silica scales. In fact, while the diffusion of interstitial molecular oxygen seems confirmed as the dominant oxygen transport process within amorphous silica, what happens when oxygen atoms are the major oxidant agent arriving at the surface does not find a shared explanation yet. Instead, the multilayer oxide scales formed upon oxidation tests in conventional furnaces and in simulated atmospheric re-entry conditions3,9,14,17,27 show however similarities. In the present case, an heat treatment at 1450 C of 20 h in flowing air at ambient pressure revealed a rather similar layered configuration.22 The oxidation of ZrB2 -SiC in air was also analysed using volatility diagrams for ZrB2 and SiC.29,30 The diagrams justified, thermodynamically, formation and stability of the outermost silica coating as well as the development of SiCdepleted regions beneath the outer oxide scales. The question whether atmospheric re-entry conditions characterised by equivalent temperatures and dwell times, induce varying thickness of the separate oxide scales is matter of future work. Other open topics connected to the consequences of the oxidation in presence of a reactive hot stream containing dissociated oxygen are currently receiving much attention. Does dissociated oxygen recombine at the surface or within the silicabased glass? Does it then diffuse inward in molecular form? In this respect, the experimental temperature-time profiles (Fig. 7) matched adequately the calculated heat flux distribution only assuming a non catalytic behavior with reference to the oxygen recombination (see specific comments in Section 4.5). It should be pointed out that the present tests have been carried out at atmospheric pressure conditions. Experimental and theoretical works on the catalytic activity of silica-based materials under simulated re-entry conditions31,32 have shown that the catalytic atomic recombination coefficients, for a constant temperature, are decreasing functions of the pressure. Therefore the catalytic properties of the material, in respect to the recombination of oxygen atoms, may be larger at lower pressures, as found for instance in arc-jet experiments with ZrB2 /SiC and HfB2 /SiC ceramic materials.2 It should also be emphasised that the co-existence of crystalline (i.e. zirconia) and amorphous (i.e. silica) phases constituting the multiphase oxide scales makes the establishment of the rate-limiting transport extremely difficult. In contrast to the amorphous silica behavior, the monoclinic crystal structure of zirconia for instance favors oxygen incorporation and diffusion in ionic rather than in molecular form. This suggest that the dissociated oxygen in the reactive gas mixture enhances oxidation since its incorporation into the zirconia scale would not require an endothermic dissociation reaction. To conclude, although the  Fig. 12. Cross-section of the ZrB2 -SiC model after arc-jet testing (surface peak temperature 2400 C): a general view (a) and details of heavily altered zones (b     and c) are shown, SEM nicrographs.     Limited applicability is however expected when the multilayer oxide scales stop mitigating the influence of high pressure inner gaseous by-products like SiO, B2O3 and CO that find ways to evolve outside disruptively. In this respect, an additional arc-jet test was carried out in the same environment but at extremely high heat fluxes (higher than 10 MW/m2 ). The temperature climbed up so rapidly to about 2400 C within 20 s that the tested model was unable to offer a valuable resistance to the aero-thermal load applied, and a stable configuration was never established. A large activity of volatile products due to the violent jump in temperature occurred in proximity of the hottest zones, leading to micro/macro-spallation of the model surface (Fig. 12a). Due to the multiform variety of the altered regions, a comprehensive description resulted an hardly solvable task. For sake of simplicity, examples of some intriguing features are given in Fig. 12b and c. The extent of the damage in these conditions is to be considered not acceptable for the foreseen demand of reliability and re-usability.  4.4. Role and influence of dissociated oxygen  The present work has been inserted into the experimental effort that has recently begun to address the question if oxidation of SiC-containing metal borides based materials in dissociated oxygen, the primary oxidant in a re-entry environment, proceeds more rapidly than that by molecular oxygen. Preliminary indications seem confirmatory.4 However, aspects like the characteristics of the oxide scales produced by oxygen atoms, or the passive-to-active oxidation boundaries of SiC remain basically unexplored. Atomic oxygen for instance is reported to accelerate oxidation rates in a variety of metals, ceramics and semiconductors.4 A series of experiments by Balat et al. showed that the temperature-pressure boundary between passive-to-active oxidation of SiC can be shifted significantly if atomic oxygen is present.25 Rogers and co-workers, studying the oxidation behavior of UHTCs and their constituents at high  \\x0c', '4804  F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805     atmospheric re-entry conditions using arc-jet testing, with an average specific total enthalpy (H) in the order of 10-20 MJ/kg. The model’s surface reached a peak value of 1930 C within 1 min for H approaching 20 MJ/kg. Elemental mapping by SEMEDS of the cross-section after exposure showed a rather compact scale of zirconia (20 \\u242em thick) underlying a thin silica-based smooth coating. A partially SiC-depleted region underneath the zirconia scale, a few microns thick, was also seen. Formation of a steady-state external multiphase oxide scales and high thermal conductivity were indicated and therefore discussed as the favorable factors which provided a considerable potential for this material in sharp leading edge applications. Numerical calculations, which simulated the chemical non-equilibrium flow around the hemispheric model correlated well with the experimental results only assuming a very low catalytic surface behavior (with reference to the oxygen recombination). Additional work needs to address the key question if this class of materials oxidises in dissociated oxygen more rapidly than in air-atmosphere. Moreover, correlate dimensional stabilityoxidation rates-active oxidation of SiC is of priority relevance.  Acknowledgements  The authors would like to thank D. Dalle Fabbriche (ISTEC) and P. Loquace (DIAS) for their technical assistance.  References  1. Upadhya, K., Yang, J.-M. and Hoffman, W., Materials for ultrahigh temperature structural applications. Am. Ceram. Soc. Bull., 1997, 76(12), 51-56.  2. Marschall, J., Chamberlain, A., Crunkleton, D. and Rogers, B., Catalytic  atom recombination on ZrB2 /SiC and HfB2 /SiC ultrahigh-temperature ceramic composites. J. Spacecraft Rockets, 2004, 41(4), 576-581.  3. Gasch, M., Ellerby, D., Irby, E., Beckman, S., Gusman, M. and Johnson,  S., Processing, properties and arc-jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics. J. Mater. Sci., 2004, 39, 5925-5937.  4. Bongiorno, A., Forst, C. J., Kalia, R. K., Li, J., Marschall, J., Nakano, A. et  al., A perspective on modeling materials in extreme environments: oxidation of ultrahigh-temperature ceramics. MRS Bull., 2006, 31, 410-418.  5. Richet, N., Lespade, P., Goursat, P. and Laborde, E., Oxidation resistance  of HfB2 -SiC coatings for protection of carbon fiber based composites. Key Eng. Mater., 2004, 264-268(TTP), 1047-1050.  6.  Janowski, R., Tauche, M., Scheper, M., Monti, R. and Savino, R., Space  plane: a new way for atmospheric reentry. In Proceedings of  the 1st Inter national ARA Days, Atmospheric Reentry Systems, Missions and Vehicles.  Session 15-System Design, 2006.  7. Monti, R., De Stefano, M. and Savino, R., Thermal shielding of a reentry  vehicle by ultra high temperature ceramic materials. J. Thermophys. Heat Transfer, 2006, 20(3), 500-506.  8. Cutler, R. A., Engineering properties of borides. In Ceramics and Glasses,  Engineered Materials Handbook, Vol 4, ed. S. J. Schneider. ASM Interna tional, Materials Park, OH, 1992.  9. Metcalfe, A. G., Elsner, N. B., Allen, D. T., Wuchina, E., Opeka, M. and  Opila, E., Oxidation of hafnium diboride. In High Temperature Corrosion  and Materials Chemistry: Per Kofstad Memorial Symposium, Electrochem ical Society Proceedings, Vol 99-38, 1999, pp. 489-501.  10. Monteverde, F., Beneficial effects of an ultra-fine ␣-SiC incorporation on  the sinterability and mechanical properties of ZrB2 . Appl. Phys. A, 2006, 82, 329-337.  11. Chamberlain, A. L., Fahrenholtz, W. G. and Hilmas, G. E., High-strength 87(6),  zirconium diboride-based  J. Am. Ceram.  ceramics.  Soc.,  2004,  1170-1172.  Fig. 13. Experimental results and numerical solutions corresponding to the dif ferent assumption of fully catalytic (FC) and non catalytic (NC) wall, 14.8 MJ/kg  of specific total enthalpy.  ZrB2 -SiC composite already displayed encouraging technical merits to be a valuable candidate in aerospace applications, many facets of the resistance to oxidation in re-entry conditions need further confirmations in order to extend this technological challenge for repeated safe uses.  4.5. Experimental-numerical correlation of aerothermal heating  Numerical computations have been carried out, using the model described in Section 3, under two different assumptions about the catalytic properties of the specimen surface with reference to recombination of atomic oxygen. The simulations refer to the plasma torch test characterised by a specific total enthalpy of 14.8 MJ/kg (whose results are shown in curve 2 of Fig. 7a) assuming the two extreme situations of a non catalytic wall and a fully catalytic wall.33 Based on the computed heat flux distributions, a thermal analysis has been carried out in both cases to evaluate the corresponding time profiles of the temperature. Fig. 13 shows the results computed under the two assumptions of fully catalytic and non catalytic wall, and the experimental profile obtained with the pyrometer. Fig. 13 shows that the experimental results match quite well the numerical ones corresponding to the non catalytic wall condition. This points out that the material herein tested exhibits, at very high temperatures, a non catalytic behavior. This behavior can be explained by the formation of a silica surface thin layer (Figs. 9-11) which is known to possess very low catalytic recombination behavior.33,34 The presence of such a surface layer also justifies the high values of the surface emissivity, according to the literature data.35  5. Summary and future work  An ultra-high temperature dense ZrB2 -SiC ceramic was produced by hot-pressing. Hemispheric ZrB2-SiC models (curvature radius 7.5 mm) were exposed to ground simulated  \\x0c', 'F. Monteverde, R. Savino / Journal of the European Ceramic Society 27 (2007) 4797-4805  4805  12. Pastor, H., Metallic borides: preparation of solid borides—sintering method  25. Balat, M.  and Berjoan, R., Oxidation of  sintered silicon carbide under  and properties of solid bodies. In Boron and Refractory Borides, ed. V. I.  Matkovich. Springer Verlag, New York, 1977.  13. Tripp, W., Davis, H. and Graham, H., Effect of an SiC addition on the oxidation of ZrB2 . Am. Ceram. Soc. Bull., 1973, 52(8), 612-616. 14. Chamberlain, A., Fahrenholtz, W., Hilmas, G. and Ellerby, D., Oxidation  of ZrB2 -SiC ceramics under atmospheric and re-entry conditions. Refract. Appl. Trans., 2005, 1(2), 1-8.  15. Monteverde, F. and Bellosi, A., Oxidation of ZrB2 based ceramics in air. J. Electrochem. Soc., 2003, 150(11), B552L B559.  16. Fahrenholtz, W., The ZrB2 volatility diagram. J. Am. Ceram. Soc., 2005, 88(12), 3509-3512.  microwave-induced CO2 plasma at high temperature: active-passive transition. Appl. Surf. Sci., 2000, 161, 434-442.  26. Hald, H., Operational  limits for re-usable space transporation systems due  to physical boundaries of C/SiC materials. Aerospace Sci. Technol., 2003, 7, 551-559.  27. Bull, D., Rasky, D. J. and Karika, J. C., Stability characterization of diboride  composites under high velocity atmospheric flight conditions. In Proceed ings of  the 24th International SAMPE Technical Conference, 2002, pp.  T1092-T1116.  28. Rogers, B., Song, Z., Marschall, J., Queralt `o, N. and Zorman, C. A., High  Temperature Corrosion and Materials Chemistry, Vol 268, ed. V. E. Opila.  17. Rezaie, A., Fahrenholtz, W.  and Hilmas, G., Oxidation  of  zirconium  The Electrochemical Society, Pennington, 2004.     diboride-silicon carbide at 1500 C at low partial pressure of oxygen. J. Am. Ceram. Soc., 2006, 89(10), 3240-3245.  18. Park, C., Nonequilibrium Hypersonic Aerothermodynamics. John Wiley &  Sons, 1990.  19. Park, C., Review of chemical-kinetic problems of future NASA missions. I. Earth Entries. J. Thermophys. Heat Transfer, 1993, 7(3), 385-398.  20. Park, C., Howe, J. T., Jaffe, R. L. and Chandler, G. V., Review of chemical kinetic problems of future NASA missions. II. Mars entries. J. Thermophys. Heat Transfer, 1994, 8(1), 9-23.  21. Park, C., Jaffe, R. L. and Partridge, H., Chemical-kinetic parameters of hyperbolic earth entry. J. Thermophys. Heat Transfer, 2001, 15(1), 76-90.  22. Monteverde, F. and Scatteia, L., Resistance to thermal  shock resistance  and to oxidation of metal diborides-SiC ceramics designed for aerospace application. J. Am. Ceram. Soc., 2007, 90, 1130-1138.  23. Thomas, D., Design  and  analysis  of UHTC leading  edge  attachment.  NASA/CR - 2002-211505, 2002.  24. De Filippis, F., Savino, R. and Martucci, A., Numerical-experimental cor relation of  stagnation point heat flux in high enthalpy hypersonic wind  29. Fahrenholtz, W., Thermodynamic analysis of ZrB2 -SiC oxidation: formation of a SiC-depleted region. J. Am. Ceram. Soc., 2007, 90(1), 143-148.  30. Wang, C. R., Yang, J.-M. and Hoffman, W., Thermal stability of refractory carbide/boride composites. Mater. Chem. Phys., 2002, 74, 272.  31. Kolesnikov, A. F., Gordeev, A. N., Vasil’evskii, S. A. and Verant, J. L.,  Predicting catalytic properties of SiC material for the Pre-X vehicle reentry  conditions. In Proceedings of the First European Conference for Aerospace  Sciences (EUCASS), 2005.  32. Kolesnikov, A. F., Yakushin, M.  I., Pershin,  I. S., Vasil’evskii, S. A.,  Chaot, O., Vancrayenest, B. et al., Comparative study of surface catalyt icity under subsonic air  test conditions.  In Proceedings of  the 4th Europ.  Symp. Aerothermodynamics  for Space Applications, 2001, pp. 481-486,  ESA SP-487.  33. Greaves, J. C. and Linnett, J. W., Recombination of atoms at surfaces. Part 5. Oxygen atoms at oxide surfaces. Trans. Faraday Soc., 1959, 55, 1346.  34. Dickens, P. G. and Sutcliffe, M. B., Recombination of oxygen atoms on oxide  surfaces. Part 1. Activation energies of recombination. Trans. Faraday Soc., 1964, 60, 1272.  tunnel. In Proceedings of  the 13th AIAA/CIRA International Space Plane  35. Wolfe, W. L. and Zissis, G.  J., The Infrared Handbook. Environmental  and Hypersonic Systems and Technologies Conference, 2005.  Research Institute of Michigan, 1978.  \\x0c']"
},{
  "_id": 247,
  "PDF": "Strength of hot pressed ZrB2–SiC composite after exposure to high temperatures (1000–1700°C).pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 32 (2012) 4455-4467  Strength of hot pressed ZrB2-SiC composite after exposure to high temperatures (1000-1700 C)     Manish Patel a,b,∗  , J. Janardhan Reddy a , V.V. Bhanu Prasad a , Vikram Jayaram b  a Defence Metallurgical Research Laboratory, Hyderabad 500058, India b Department of Materials Engineering, Indian Institute of Science, Bangalore 560012, India  Received 29 May 2012; received in revised form 22 June 2012; accepted 28 June 2012  Available online 31 July 2012  Abstract  Residual strength (room  temperature strength after exposure  in air at high  temperatures) of hot pressed ZrB2 -SiC composites was evaluated as function of SiC contents (10-30 vol%) as well as exposure temperatures for 5 h (1000-1700 C). Multilayer oxide scale structures were found after exposures. The composition and  thickness of  these multilayered oxide scale structure was dependent on exposure  temperature and SiC contents in composites. After exposure to 1000 C for 5 h, the residual strength of ZrB2 -SiC composites improved by nearly 60% compared to the as-hot pressed composites with 20 and 30 vol% SiC. On the other hand, the residual strength of these composites remained unchanged after 1500 C for 5 h. A drastic degradation in residual strength was observed in composites with 20 and 30 vol% SiC after exposure to 1700 C for 5 h in ZrB2 -SiC. An attempt was made to correlate the microstructural changes and oxide scales with residual strength with respect to variation in SiC content and temperature of expsoure. © 2012 Elsevier Ltd. All rights reserved.              Keywords: A. Hot pressing; D. Borides; Ultra High Temperature Ceramics; Oxidation; Flexural strength  1.   Introduction     In  the  recent past, considerable work has been carried out in  the field of ultra high  temperatures ceramics, especially on ZrB2 based ceramics and composites, due to interest in technologies associated with hypersonic flight, atmospheric re-entry, and rocket propulsion. Materials  in such applications must possess high melting temperature (3500 C) along with high electrical and thermal conductivities, good oxidation and chemical attack resistance.1-3 The high temperature oxidation resistance of pure ZrB2 is not sufficient for the aggressive environments associated with the above applications. Addition of SiC as reinforcement in ZrB2 is known to improve the oxidation resistance by enabling the  formation of protective borosilicate glass  layer4-5 which offers a higher viscosity, higher boiling point and  lower vapor pressure  than B2O3 ,  thereby providing more efficient oxidation protection.6 It was found that the oxide scale of ZrB2-SiC  ∗  Corresponding author at: Ceramic and Composite Group, Defence Metal lurgical Research Laboratory, Kanchanbagh, Hyderabad 58, India.  Tel.: +91 4024586824; fax: +91 4024340683.  E-mail address: patelmet@yahoo.co.uk (M. Patel).  0955-2219/$ - see front matter © 2012 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2012.06.025  composites generally had a  layered structure. Based on detail work by Zhang et al.,7 the structural changes in the oxide scale revealed:        (i)   (ii)   In the temperature range from 700 C to 1200 C, the oxide structure consists of a B2O3 -rich outer  layer, a subscale of ZrO2 that contained unoxidized SiC, and unaffected ZrB2-SiC in the substrate. In  the  temperature  range 1200-1600 C,  the outer  layer changed to a SiO2 -rich glassy oxide followed by a subscale of ZrO2 that contained some SiO2 and, finally, unaffected ZrB2-SiC. (iii) At temperatures above 1600 C, an additional SiC-depleted layer forms before one reaches  the unaffected ZrB2-SiC. The SiC-depleted layer develops due to the change of oxidation kinetics of SiC from passive to active.8        For re-usable applications, it is essential to know the effect of these oxide layer structures on residual strength of the ZrB2-SiC composites at room temperature after exposure to high temperatures. Few reports are available on residual strength of ZrB2 -SiC composite after thermal exposure at high temperatures.9-11 Guo        \\x0c', '4456   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467        and Zhang9 have shown that the flexural strength of ZrB2 with 10 and 30 vol% SiC increased sharply after exposure at 1500 C for 30 min and then gradually decreased for longer times. Recently, it was shown that the oxide scale structure has a great influence on the improvement of strength of ZrB2-SiC-graphite composC.12 Previous work by authors shows ite after exposure to 1100 the effect of B4C content on  residual strength of hot pressed ZrB2 after exposure  to high  temperatures.13 It was found  that the residual strength of hot pressed ZrB2 with 0.5 and 1 wt% of B4C dramatically  increased after exposure  to 1000 C for 5 h. In contrast,  larger volume fractions of B4C (3 and 5%) did not lead to any improvement. Previous studies on residual strength of ZrB2 -SiC composites have been limited to specific exposure temperature.9-12 But there is a need to measure residual strength after exposure to a range of  temperatures  since  the oxide  layer  structure changes with temperatures. The present work seeks  to address  this gap by measuring the dependence of residual bend strength on thermal exposure under oxidizing conditions at different  temperatures and times.     2. Experimental details  (d50         (d50  7   Commercially  available ZrB2 powders   1.5-3  \\u242em), supplied by H.C. Strack, Germany, SiC powders (d50   1.5  \\u242em), supplied  by  H.  C.  Strack,  Germany  and  B4C  powders \\u242em), supplied by Electro Abrasive Corporation, USA were used as raw materials. ZrB2-SiC composites with 10, 20 and 30 vol% SiC were prepared. All these composites had small amount of B4C (1 wt% of ZrB2 contents) as additive. Powders were mixed for 24 h in a polythene bottle with alumina balls and then hot-pressed in graphite dies at 2000 C with a heating rate of 15 C/min under a uni-axial pressure of 25 MPa for 1 h. Hot pressed composites were named as Z10S, Z20S and Z30S for 10, 20 and 30 vol% of SiC. The bulk density was measured using the Archimedes displacement method with water as the immersing medium and the relative density was calculated with respect to theoretical density. The theoretical density was estimated using rule-of-mixtures calculations based on  the  initial composition of  the samples. The polished surface of  the hot pressed sample was  revealed by SEM without etching. Flexural  strength was measured by the three-point bending method with a span of −1 using an Instron 40 mm and a cross-head speed of 0.5 mm min 5500R universal  testing machine. The size of  the samples was 3   50 mm. Bend test samples were prepared from diamond ground, hot pressed discs by EDM wire cutting. For  residual strength measurement at room temperature, the bend test specimens of hot pressed ZrB2-SiC  composite were  exposed  to 1000 C, 1500 C and 1700 C in an air furnace. Specimens were supported on Al2O3 fixtures for 1000 C and 1500 C and ZrO2 fixture  for 1700 C. Specimens were heated  to  the designated temperature at a heating  rate of 10 C/min. Specimens were held  for 5 h at  the designated  temperature and  furnace cooled to room  temperature. Residual strength of exposed specimens was measured at room temperature as described previously. Xray diffraction of the oxide layer was performed using a Philips  ×  ×   4                        X-ray diffractometer with copper radiation as the X-ray source. The  relative composition of monoclinic ZrO2 and  tetragonal ZrO2 was calculated  from XRD pattern using  the polymorph method14 which uses the (1 0 1) peak of tetragonal (t) ZrO2 and (−1 1 1) and  (1 1 1) peaks of monoclinic  (m) ZrO2 as shown below:  Vt-ZrO2  =   1   −  1   +  1.311(1    0.311(1    Xt-ZrO2 )  Xt-ZrO2 )  −  −  =  Xt-ZrO2  +  I(1 0 1) (t)   I( ¯1 1 1) (m)   +  I(1 0 1) (t)    I(1 1 1) (m)  I(m) are  the peak (m) ZrO2 ,  respec where Xm is  the  intensity  ratio and  I(t) and  intensities of  tetragonal  (t) and monoclinic  tively and V is the volume fraction. The weight of  specimens before  and  after  exposure was measured  for calculation of weight change due  to oxidation. The oxide  layer  thickness was measured  from  fracture crosssections. Fracture surfaces of the specimens were examined by scanning electron microscopy. Micro-hardness was measured by Vickers  indentation  tests using a  load of 500 g and a dwell time of 10 s. The  indentation fracture  toughness was measured from direct crack measurement from following equation15 :  (cid:2)  (cid:3)1/2(cid:2)  (cid:3)3/2  KC =   0.016  E  H  P  c  where E is the elastic modulus, H is the hardness, P is the load and c is the crack length from the center of indent. Average microhardness and indentation fracture toughness was calculated from five indents. Elastic modulus was measured by pulse-echo technique, using following relation:  (cid:4)  −  (1   (cid:5)  +   ν)   2ν)(1   −   ν)  (1   =  E    ρV 2  L  where ρ  is density of the sample, VL is longitudinal velocity of ultrasonic waves and ν  is  the Poisson ratio. The Poisson ratio was assumed to be 0.16 of ZrB2-SiC composites.1 Longitudinal ultrasonic velocity was measured using a 10 MHz transducer.  3. Results  3.1. As hot pressed composite  Fig. 1  shows  the microstructure of hot pressed ZrB2 -SiC composites. The light phase is ZrB2 whereas the dark phase consists of SiC. Near-theoretical density of composite was achieved by hot pressing  for ZrB2-SiC composites with different content of SiC (Table 1). The average grain size of ZrB2 decreased with  increase of SiC contents. The average grain size of ZrB2 \\u242em.  in Z10S samples was  found  to be 4.3   1.4  It decreased marginally to 3.9   0.9  \\u242em in Z30S samples suggesting a small role of SiC  in restricting  the coarsening of ZrB2 . The average micro-hardness and fracture toughness of ZrB2-SiC composites increased with SiC content. The average micro-hardness of Z10S was found to be 18.0   0.9 GPa and increased to 24.4   0.6 GPa for Z30S, reflecting  the higher hardness of SiC (27 GPa) compared  to ZrB2 (15-18 GPa). The flexural  strength of  as hot  ±  ±  ±  ±  \\x0c', 'M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   4457  Fig. 1. Microstructure (BSE image) of hot pressed ZrB2 -SiC composites, (a) Z10S, (b) Z20S and (c) Z30S. Bright and dark phases are ZrB2 and SiC, respectively.  Fig. 2.   (a) Variation of flexural strength and (b) fracture surface of as hot pressed ZrB2 -SiC composites, showing both transgranular and intergranular mixed type fracture.  ±  ±  pressed ZrB10SiC composite was found  to be 393   114 MPa, which increased to 487   68 MPa for Z20S. Further, the flexural strength was decreased  to 425   53 MPa  for Z30S  composite  (Fig. 2a). The mode of  fracture of hot pressed ZrB2-SiC composite was observed as mixed type (both intergranular and intragranular) brittle fracture (Fig. 2b).  ±  3.2. Oxidation  3.2.1. Weight gain  Fig. 3 shows  the variation of weight gain per unit area of different ZrB2 -SiC composite after exposure  to different  temperatures  for 5 h. Plot clearly shows  that  the weight gain was considerable increased with exposure temperature from 1000 C to 1500 C. Further  increasing  the exposure  to 1700 C, more weight gain was noticed. Weight gain decreased with increase in SiC contents in ZrB2-SiC composites. With exposure to 1000 C for 5 h, the relative change of weight gain with SiC content was              2  m m  /  g  m  ,  n  i  a  G  t  h g  i  e  W  0.35  0.30  0.25  0.20  0.15  0.10  0.05  0.00  After e exposu re at  100 000C for 5 hour rs After e exposu re at  150 000C for 5 hour rs After e exposu re at  170 000C for 5 hour rs  5  10  15  20  25  30  35  SiC content, Vol %  Fig. 3. Weight gain of hot pressed ZrB2 -SiC composites with SiC contents after exposure to 1000 C, 1500 C and 1700 C for 5 h.           Table 1  Microstructure and mechanical properties of hot pressed ZrB2 -SiC composite.  Sample ID   SiC content (vol%)   Relative density (%)   Grain size (\\u242em)   Elastic modulus (GPa)   Micro-hardness (GPa)   Fracture toughness (MPa m1/2 )  Z10S  Z20S  Z30S  10  20   30   99.8   99.7   97.5   4.29   4.02   3.88    1.42    1.1   ± ± ±   0.89   ±   16   ±   12   500   506   487   17.98   22.13   24.36   ± ± ±   0.9    1.1    0.6   3.76   4.15   4.44   ± ± ±   0.34   0.83   0.48                \\x0c', '4458   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467  Fig. 4. Oxide scale morphologies from after exposure to 1000     C for 5 h of Z10S. (a) Photograph of exposed specimens, (b) the BSE image of top horizontal surface,  showing presence of both ZrO2 (bright phase) and glassy B2O3 (dark phase), (c)  amount of B2O3 .  very low. The weight gain after exposure to 1500 C is related to the relative amounts of ZrO2 and SiO2 phase in the oxide scale.        3.2.2. Oxide scale morphology 3.2.2.1. Top surface view. Top surface of ZrB2-SiC composite     after exposure  to 1000 C for 5 h showed a clear difference  in color between the top horizontal surface and vertical surface (as indicated  in Fig. 4a). Horizontal and vertical surface positions in Fig. 4a were similar  to position  in furnace during exposure. The color of  top surface after exposure  to 1000 C for 5 h was gray but top horizontal surface was darker than top vertical surface. Difference  in color of  top surface was directly matched with difference in morphology of the top surfaces. Fig. 4b and c shows back scattered electron SEM images of the top horizontal surface and vertical surface,  respectively. The bright phase  in back scattered  image of  top surface was crystalline ZrO2 and dark phase was glassy phase. The presence of crystalline ZrO2 in  the oxide scale was confirmed by XRD as monoclinic ZrO2 (Fig. 5a). From Fig. 4b and c,  it was found  that  the amount of glassy phase present was more in the top horizontal surfaces as compared to the top vertical surfaces. On other hand, the distribution of glassy phase on top horizontal surface varied with SiC content after exposure to 1000 C for 5 h. The agglomeration of glassy phase was found to be more pronounced on top surface of Z20S than in Z10S and Z30S, as shown in Figs. 4b and 6a and b. The large agglomeration of glassy phase on top surface of Z20S was well correlated with the oxide layer thickness as the oxide layer  thickness was  the  lowest for Z20S. On  the  top surface of Z30S, cracks were found  in glassy regions  throughout  the surface. The  inset  in Fig. 6b shows  the crack  formed on  the  top surface of Z30S. Some of  the darker phase on  top surface as appeared  to shape of SiC  in un-oxidized composites (Fig. 6c) whereas the EDX analysis on these phases revealed the presence of SiO2 (Fig. 6d). The top surface after exposure to 1500 C for 5 h was blackish in color with glowing luster having white spots, distributed on the surfaces (Fig. 7a). The numbers of these white spots decreased        the BSE   image of   top vertical surface, showing presence of ZrO2 with very   low  with increase in SiC content, as shown in Fig. 7b and c. Fig. 7b and c shows the morphology of these white spots. It looks likes agglomerates of small particles. The BSE image of top surface differentiated  into  three different contrasts (Fig. 8a). Based on these atomic contrasts,  the darkest  region  is B2O3 rich glass whereas  the brightest region  is ZrO2 . The reaming contrast  in Fig. 8a  is from SiO2 rich glass, which was shown  the glowing luster of  top  surfaces. The presence of crystalline ZrO2 was identified by XRD (Fig. 5b). The crystalline ZrO2 on top surface after exposure  to 1500 C  for 5 h was shown  to a mixture of monoclinic and tetragonal phases (Fig. 5b). The relative content of tetragonal ZrO2 increased with increase in SiC contents. The volume fraction of tetragonal ZrO2 was found to be 4% and 18% for Z10S and Z30S (Table 2), respectively. A closer view of these white spots, cracks between  the ZrO2 particles  (Fig. 8b) and     150  100  50  0  200  100  400  300  200  100  0  y  t  i  s  n  e  t  n  I  )  1 1 1  )  1 1 1  (  ( )  m )m  1 1 0 0 1 1  ( ( t t  m-ZrO2 & & t-ZrO2  m-ZrO2  & t-ZrO2  m-ZrO  2  c  b  a  20  30  40  50  60  70  80  90  2θ     (b) 1500  C and (c) 1700  Fig. 5. XRD pattern of top surface of Z10S after exposure for 5 h to (a) 1000 C, −1 1) peaks of monoclinic ZrO2 and (1 0 1) peak of tetragonal ZrO2 are only indexed since these peaks were used in calculation of relative composition of tetragonal ZrO2 .  C. (1 1 1) and (1         \\x0c', 'M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   4459  Table 2  Relative composition of tetragonal and monoclinic ZrO2 on top surface of oxide scale formed on ZrB2 -SiC composites.  Sample ID   Z10S  Z20S  Z30S     1500  C for 5 h      1700  C for 5 h  4%   11%   18%   10%  20%  18%        dendrites of ZrO2 (Fig. 8c) were observed. These ZrO2 particles and dendrites are called as secondary ZrO2 , which are different from ZrO2 formed directly due to oxidation of ZrB2 . Similar to exposure at 1500 C, the top surface after exposure at 1700 C was shown to be a mixture of monoclinic and tetragonal ZrO2 by XRD. The  relative content of  tetragonal ZrO2 increased with  increase of SiC contents. The volume  fraction of  tetragonal ZrO2 was  found  to be 10% and 20%  for Z10S and Z20S (Table 2), respectively. The top surface of ZrB2-SiC composite after exposure  to 1700 C for 5 h had open (marked     Fig. 6. Oxide scale morphologies on   top surface after exposure   to 1000     C for 5 h, (a) BSE   image for Z20S, (b) BSE   image for Z30S, (c) SE   image for Z20S, (d)  EDX pattern for phase shown as arrows. The inset in (b) is magnified view, showing cracks in glassy phase.  Fig. 7. Oxide scale morphologies on top surface after exposure to 1500  SE image for Z30S, showing agglomerates of ZrO2 on glassy SiO2 .     C for 5 h. (a) Photograph of exposed specimens for Z10S, (b) SE image for Z10S and (c)  \\x0c', '4460   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467  Fig. 8. BSE   image of oxide scale morphologies on   top surface of Z20S after exposure   to 1500     C for 5 h. (a) Low magnification   image shows different atomic  contrast, (b) magnified view image ZrO2 agglomerates, showing cracks between ZrO2 particles and (c) Higher magnification image of ZrO2 agglomerates shows the dendrite form of ZrO2 .  Fig. 9. Oxide scale morphologies on top surface after exposure to 1700     C for 5 h. (a) Photograph of exposed specimens, bubbles are shown as red circles, SE image  from inside of bubbles for (b) Z10S and (c) Z30S. Images show particulate ZrO2 and dendrite ZrO2 in bubbles of Z10S and Z30S, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)     as  red circles  in Fig. 9a) and closed bubbles. With  increasing SiC content after exposure  to 1700 C for 5 h,  the size of bubbles on  top decreased  though  their number density  increased. ZrO2 was found on SiO2 rich oxide  in  the  inside of open bubbles, as shown  in Fig. 10. The morphology of ZrO2 particles was different for ZrB2 -SiC composite with different SiC contents. Particulate ZrO2 was found  in  the bubbles of Z10S and  Z20S (Fig. 9a). The morphology of ZrO2 in the bubbles of Z30S was dendrites (Fig. 9b). Similar  to morphology of ZrO2 in  the bubbles, the morphology of ZrO2 on remaining top surface was different for different SiC contents, as shown in Fig. 10. Particulate morphology of ZrO2 was  found on  top surface of Z10S and Z20S  (Fig. 10a) whereas well decorated ZrO2 dendrites structures was noticed on top surface of Z30S (Fig. 10b).  Fig. 10. SE image of oxide scale morphologies on top surface outside the bubbles after exposure to 1700  ZrO2 dendrites on oxide surface of Z30S.     C for 5 h (a) ZrB20 SiC and (b) Z30S. It shows network of  \\x0c', 'M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   4461  Fig. 11.   (a) Cross-sectional view of oxide scale of Z10S (BSE image) and (b) oxide layer thickness variation with SiC contents after exposure to 1000     C for 5 h.        ±  3.2.2.2. Cross-sectional view. A   single  layer oxide  structure was observed  in cross section of oxidized samples after exposure  to 1000 C  for 5 h  (Fig. 11a). The oxide  layer contained ZrO2 rich oxide scale with un-reacted ZrB2-SiC. The ZrO2 rich oxide  scale contained glassy B2O3 and un-reacted SiC apart from ZrO2 , as shown  in Fig. 11a. Fig. 11b shows  the variation of oxide layer thickness variation after exposure to 1000 C for 5 h for hot pressed ZrB2-SiC composite. The thickness of ZrO2 rich oxide  layer was minimum for Z20S samples (15  \\u242em). The  thickness of ZrO2 rich oxide  layer was 31.4   8.7  \\u242em and 21.2   4.4  \\u242em for Z10S and Z30S, respectively. A  two-layer structure was observed  in cross section of oxidized samples after exposure  to 1500 C  (Fig. 12a). A dense SiO2 outer layer followed by a middle layer of ZrO2 rich oxide was observed. The ZrO2 rich oxide scale had glassy SiO2 (dark contrast  in ZrO2 rich oxide  region)  in between ZrO2 particles. The formation of distinct SiO2 rich and ZrO2 rich oxide layer was also  revealed by EDX elemental mapping of Zr, Si and O, as shown  in Fig. 13. The  thickness of dense SiO2 rich oxide  layer  after  exposure  to 1500 C  for 5 h was  found  to be 24   10.6  for Z10S. The  thickness of SiO2 rich oxide layer  increased  to 64   15  \\u242em for Z20S and  then decreased  to  \\u242em   ±  ±  ±  ±   2         ±        ±  ±  ±  47.3   10.3  \\u242em for Z30S. But the trend in the ZrO2 layer was different (Fig. 12b). The thickness of ZrO2 rich oxide scale for Z10S was 140   45  \\u242em. It decreased to 50   18  \\u242em for Z20S and then further decreased with SiC content to 34   12  \\u242em for Z30S. Similar  to exposure at 1500 C, a  two-layer  structure was observed  in cross section of oxidized samples after exposure to 1700 C  for 5 h  (Fig. 14a). The outer  layer was SiO2 rich oxide  layer followed by  thick ZrO2 rich  layer. The ZrO2 particles  in ZrO2 rich  layer were oriented  toward  the outer surface, as shown  in Fig. 14a. The  thickness of SiO2 rich oxide outer layer was found  to be 40.5   9.4  \\u242em for Z10S and 46  \\u242em for Z20S. For Z30S, the thickness of SiO2 rich oxide outer layer (128   23  \\u242em) was thicker than Z10S and Z20S. The thickness of ZrO2 rich oxide layer was 470   45  \\u242em for Z10S, which was comparatively larger than thickness of ZrO2 rich oxide layer for Z20S (277   20  \\u242em) and Z30S (101   17  \\u242em). ZrO2 particles in SiO2 rich outer  layer were different from  the ZrO2 particles in ZrO2 rich oxide layer. The ZrO2 particles in ZrO2 rich oxide layer are primary ZrO2 which formed directly from oxidation of ZrB2 . The size of primary ZrO2 particles in ZrO2 rich oxide layer was found to be different at different position from inner boundary  to outer boundary, as  indicated as arrow ZrO2 rich oxide  ±  ±  ±  ±  ±  ±   6   Fig. 12.   (a) Cross-sectional view of oxide scale of Z20S (BSE   image) and (b) oxide   layer   thickness variation (both ZrO2 rich   layer and SiO2 rich   layer) with SiC  contents after exposure to 1500  C for 5 h.     \\x0c', '4462   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467  Fig. 13. EDX elemental map for Zr, Si and O of cross-sectional image of oxide scale structure of Z20S after exposure to 1500     C for 5 h.  ±  ±  layer. The size of primary ZrO2 particles in Z10S is almost constant (19.6   4.7  \\u242em for middle of boundary and 21  \\u242em for outer boundary)  through out  the  thickness of ZrO2 rich oxide layer, except near the boundary between ZrO2 rich oxide layer and un-reacted ZrB2-SiC composite (Fig. 15c). Near the boundary between ZrO2 rich oxide  layer and un-reacted ZrB2-SiC composite  in Z10S  is  fully sintered  (Fig. 15a). Sintered zone of ZrO2 rich oxide  layer showed  the presence of Zr and O  in EDX analysis (Fig. 15d). Whereas, near the boundary between ZrO2 rich oxide  layer and un-reacted ZrB2-SiC composite  in Z20S and Z30S, ZrO2 particles are separated by glassy phase (Fig. 15b). The size of primary ZrO2 particles at inner boundary were larger (18.6  \\u242em) than middle (9.6   3.2  \\u242em) and outer boundary (8.3   3.7  \\u242em) for Z20S. Similar  to Z20S,  the ZrO2  ±  ±  ±   7    8    1   ±  ± ±  particles  in ZrO2 rich oxide  layer of Z30S were  larger at  inner boundary (12.7   5.3  \\u242em) than middle (4.5   1.6  \\u242em) and outer boundary (4.7  \\u242em). Both the ZrO2 particles at inner boundary and outer boundary decreased with increase in SiC content. Smaller ZrO2 particles are due to presence of SiO2 in between the ZrO2 particles. As SiC content increased, the amount of ZrO2 in ZrO2 rich oxide layer decreased. It was found to 90 vol% of ZrO2 in ZrO2 rich oxide layer for Z10S where as only 59 vol% ZrO2 in ZrO2 rich oxide  layer for Z30S (Fig. 14b), remaining volume occupied by SiO2 and pores. The pores were more  in Z10S than Z20S and Z30S. Also, the amount of SiO2 increased from  inner boundary  to outer boundary. The ZrO2 particles  in SiO2 rich outer layer were secondary ZrO2 , which formed from precipitation borosilicate liquid during solidification.  Fig. 14.   contents after exposure to 1700  (a) Cross-sectional view of oxide scale of Z20S (BSE   thickness variation (both ZrO2 rich  layer and SiO2 rich  C for 5 h. The volume fraction of ZrO2 content in ZrO2 rich oxide layer is shown for each composites.  image) and (b) oxide   layer      layer) with SiC  \\x0c', 'M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   4463  Fig. 15. SE   after exposure to 1700  of arrow in (a).  image of cross-sectional view of oxide scale at higher magnification near   the boundary between ZrO2 rich  layer and un-reacted ZrB2 -SiC composite C for 5 h (a) Z10S, (b) Z20S, panel (c) shows the variation of ZrO2 particles size through out the ZrO2 oxide scale thickness. (d) EDX pattern                       ±  ±  ±  ±  3.2.2.3. Residual strength. Fig. 16 shows   the residual flexural strength of hot pressed ZrB2-SiC composite at  room  temperature after exposure  to 1000 C, 1500 C and 1700 C for 5 h. The plot clearly shows  that  the effect of oxidation on residual strength for Z10S composite was different form Z20S and Z30S composites. For Z10S composite, the flexural strength was found to be 519   119 and 515   29 MPa after exposure  to 1000 C and 1500 C  for 5 h,  respectively. Compared  to strength of as hot pressed Z10S,  30% (Fig. 16b). Exposure to 1700 it was an  improvement  in strength value by nearly  C for 5 h, the flexural strength of Z10S composite was found to be 374   27 MPa, retaining the original strength of composite. On other hand, the flexural strength of Z20S and Z30S composites was found to be 800   47 and 674   35 MPa after exposure  to 1000 C for 5 h, respectively. This was a considerable  improvement  in strength values (60%) compared to as hot pressed condition for Z20S and Z30S composites. After exposure  to 1500 C  for 5 h,  the strength of Z20S and Z30S composites was found  to 445   65 and 414   35 MPa, respectively. Unlike the Z10S composition, Z20S and Z30S showed no significant change  in strength after exposure  to 1500 C  for 5 h. A drastic degradation  in flexural strength value of Z20S and Z30S composites after exposure  to 1700 C for 5 h was observed. This was nearly degraded to 70% to strength of hot pressed condition. The flexural strength value after exposure  to 1700 C for 5 h for Z20S and Z30S composites was found  to be 156   10 and 112   35 MPa, respectively. The flexural strength of Z30S composite was  lower  than Z20S     ±  ±  ±        ±        ±  composite after exposure  5 h.  to 1000     C, 1500     C and 1700     C for  4. Discussion  4.1. Strength of as hot pressed composite  Zhang  et  al.16 has  shown  that  the  flexural  strength  of ZrB2-SiC  composite,  prepared  by  pressureless  sintering, increases with  increases  in SiC content  in composites. Similar dependence of flexural  strength on SiC contents  is also reported for hot pressed ZrB2 -SiC composite.17 But the present result shows a different trend of flexural strength for hot pressed ZrB2 -SiC  composites  (Fig.  2a). Flexural  strength  increased by  increasing  the  SiC  content  from  10 vol%  to  20 vol%, but  decreased  for  30 vol% SiC  composite  as  compared  to 20 vol% SiC composite. From Table 1,  it was found  that grain size decreases with  increases  in SiC  contents  in ZrB2 -SiC composites. The corresponding hardness and  fracture  toughness  increased with SiC  contents,  as  the  hardness  of SiC (25-30 GPa16 )  is more  than ZrB2 (15-18 GPa13 ). The  trend of flexural strength dependence on volume fraction of SiC may be due to counter balancing the increment in strength due to grain size and residual  tensile stress generates  in ZrB2 phase during −6 / cooling after hot pressing since the CTE of ZrB2 (6.8   10 C) × 10 −6 / is higher than SiC (4.5  C).  ×        \\x0c', '4464   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467  5  10  15  20  25  30  0  100  200  300  400  500  600  700  800  900  1000  F  l  x e  u  r  a  l  s  t  r  e  g n  t  h  ,  M  P  a  SiC content, %   As  hot  press ed  After exposu re  to 10  After exposu re  to 15  After exposu re  to 17  000C 000C 000C  -100 900  1200  1500  1800  -80  -60  -40  -20  0  20  40  60  80  100  D  e  g  r  a  d  a  t  i  n o  I  m  p  r  o  e v  m  e  n  t  C  h  a  g n  e  i  n  A  e v  r  a  g  e  f  l  x e  u  r  a  l  s  t  r  e  g n  t  h  ,  %  Exposure Temperature, C   10 vol% SiC  20 vol% SiC  30 vol% SiC  No I mprovement  a  b  Fig. 16.   (a) Residual strength of ZrB2 -SiC composites after exposure to 1000 exposure temperatures.     C, 1500     C and 1700     C for 5 h, and (b) percentage change in residual strength with  4.2. Oxidation and residual strength  4.2.1. Exposure to 1000     C for 5 h  The oxidation behavior of ZrB2-SiC composites is different in terms of weight gain and oxide scale thickness after exposure to 1000 C  for 5 h. The decrease  in weight gain of ZrB2 -SiC composite with  increase  in SiC  content  shows no oxidation of SiC phase  in ZrB2 -SiC composite. The oxidation of both ZrB2 and SiC results in weight increase as both ZrO2 and SiO2 has higher molecular weight than their constituent un-oxidized phases. If both the phases oxidize, the weight gain must increase with  increase  in SiC contents  in ZrB2-SiC composite. But  the presence of SiO2 phase on SiC shows partial oxidation of SiC at 1000 C for 5 h. It has been reported that the oxidation of SiC phase  in ZrB2-SiC composites start at  temperature >1100 C due to perennial oxidation of ZrB2 .18 But the oxidation of micron C.19 So  size SiC  single phase  starts as early as at 843 it  is possible  to partially oxidize SiC  in ZrB2-SiC at 1000 C  for 5 h. Previous work by authors13 has shown  that part of B2O3 remains in the oxide scale of ZrO2 after oxidation of hot pressed ZrB2 having 1 wt% B4C after exposure  to 1000 C for 5 h. So, the glassy phase  in  the oxide scale  is B2O3 rich silicate glass having  some amount of SiO2. The binary phase diagram  for 75% of SiO2 SiO2 -B2O3 shows that B2O3 can dissolves up to  at 1000 C.5 So, oxidation of ZrB2-SiC composite at 1000 C can be represented by chemical reactions (1)-(3).                          ZrB2 +   5O2 =   2ZrO2 +   2B2O3 (l)   (1)  2SiC   +   3O2 =   2SiO2 +   2CO   (2)  2B4C   +   7O2 =   4B2O3 (l)   +   2CO   (3)  Chemical reaction (3) represents the oxidation of B4C, which is used as sintering additive whereas glassy B2O3 (l) evaporates by chemical reaction (4).  B2O3 (l)   =   B2O3 (g)   (4)  As partial oxidation of SiC occurs after exposure to 1000 C for 5 h,  formation of SiO2 is  responsible  for  lower  thickness of oxide scale of Z20S composite as compared  to Z10S. But the increased oxide scale thickness of Z30S compared to Z20S appears  to be due  to  the  formation of cracks on  the  top surface of oxide scale. Cracks formed in oxide layer in ZrB2 -SiC composite may be due to residual stresses generated by ionic diffusion, growth of ZrO2 , pressure exerted by CO (g) and volatile B2O3 (g) formed  in subsurface, volume changes during oxidation and CTE mismatch between oxide products, as  reported by Mallik et al.18 Since  the cracks were  found only on B2O3 oxide surface of Z30S, stresses generated due to ionic diffusion, growth of ZrO2 can be ruled out. On the other hand, formation of CO  (g) and content of SiO2 in glassy phase due  to partial oxidation of SiC by reaction (2) increases with increase in SiC contents. CTE of glassy phase will decrease with  increases of −6 / SiO2 content, as CTE of SiO2 is very  low  (0.5   10 C). Therefore,  the possibility of cracks  in glassy phase on  top surface of Z30S due  to stresses by CTE  is  less since glassy phase experienced more compressive stress on  top surface of Z30S. So,  it  is possible  that  the cracks on  the oxide surface of Z30S are due to stresses generated from CO (g). Due to formation of these cracks on  the  top surface of Z30S,  the oxide scale  thickness was higher for Z30S than Z20S after exposure to 1000 C for 5 h. On other hand, the variation of oxide scale thickness has shown  to direct correlation for residual strength with SiC contents. So, the improvement in strength after oxidation at 1000 C for 5 h is attributed to the healing of surface defects. Previously, authors have  shown  that, oxide  scale with minimum defects increases  the flexural  strength of hot pressed ZrB2 considerably after exposure  to 1000 C  for 5 h.13 Also, a pre-cracked ZrB2-SiC composite shows strength  improvement by healing of pre-crack at 1100 C.12 Since, Z20S and Z10S has lower the oxide  scale  thickness  than Z10S  (Fig. 11b) and also  the  top surfaces of Z20S and Z30S  (Fig. 6a and b) has shown better glassy phase distribution than Z10S (Fig. 4b). This resulted into considerable improvement in strength values (60%) for Z20S and Z30S as compared to slight improvement in strength values     ×                                                         \\x0c', 'M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   4465     (30%)  for Z10S after exposure  to 1000 C  for 5 h. Further, the strength after exposure  is  lower  for Z30S  (Fig. 16a)  than Z20S, as the oxide layer thickness of Z30S is higher than Z20S (Fig. 11b). This can also be correlated with probability of healing of flaws  is expected  to  the highest with  lowest oxide scale thickness.  4.2.2. Exposure to 1500     C for 5 h           Formation  of  two  distinct  oxide  layer  structures  (SiO2 layer  and ZrO2 layer)  in ZrB2-SiC  composite  after  oxidation  at  1500 C  in  air  is  very well  reported  by  different researchers.7,11,20 Formation of dendrites of ZrO2 in agglomerates  of ZrO2 (Fig.  8b  and  c)  on  SiO2 layer  shows  the precipitation of ZrO2 from  liquid,  similar  to  the mechanism purposed by Karlsdottir et al.5,21,22 for  transportation of ZrO2 with borosilicate glass  through convection cells. As shown by Karlsdottir et al.,23 the  formation of  less number of convection  cells  (agglomerates of ZrO2 in Fig. 7b  and  c),  as SiC contents increases in composites is due to formation SiO2 content  increases, which dissolves  less amount of ZrO2 in  liquid. Cracks  in  agglomerates of ZrO2 are  formed due  to volume expansion during phase  transformation of ZrO2 . The  stable form of ZrO2 at 1500 C  is  tetragonal, which  transform  into monoclinic at 1170 C. The tetragonal to monoclinic ZrO2 transformation during cooling is associated with a volume expansion of 3-6%,18 which  imposed a  tensile  residual stress  in glassy phase between ZrO2 . This resulted into cracks in agglomerates of ZrO2 . Thermodynamically, the oxidation of ZrB2 -SiC composites at 1500 C can be  represented by chemical  reactions (1)-(4). But, the partial pressures of O2 decreases with increase in oxide layer thickness.24 This activates the chemical reaction (5) by direct oxidation of SiC  into SiO  (g) which again converted into liquid SiO2 by chemical reaction (6) at outer surface of oxides. So, there is no loss of SiO2 as SiO (g) at 1500 C. The above reaction can be also derived from variation of weight gain and oxide  layer  thickness with SiC content. Weight gain and ZrO2 oxide layer decreased with increase in SiC, which shows that oxidation resistance of ZrB2 -SiC composites at 1500 C for 5 h improved with SiC contents.  O2 =  SiO(g)   O2 =   2SiO2 (l)   2SiO(g)    CO   SiC   +  +  +  (5)  (6)                 Residual strength of ZrB2-SiC composites after exposure to 1500 C  for 5 h  follows  the opposite  trend of variation outer SiO2 oxide  layer with SiC contents. The behavior  is similar  to residual strength variation with SiC contents after exposure  to 1000 C for 5 h. Since  the SiO2 oxide  layer  thickness of Z10S is  lower  than Z20S and Z30S,  the residual strength of Z10S  is slightly higher  than Z20S and Z30S. But  the overall strength after exposure to 1500 C for 5 h with respect to as hot pressed is  (±10%) was considered  in no  improvement zone  (Fig. 16b). No  improvement zone the experimental variation  in strength measurement, as the strength of ceramics shows statistical variation. As the published reports on residual strength of ZrB2-SiC composites after exposure  to 1400-1500 C  showed conflicting  in results. Zhang et al.9 reported  that  the residual strength              of ZrB2-20 vol% SiC  composite  is decreased  to 80% of  as processed composite after exposure at 1400 C for 0.5 h, which further  improved by 21% after exposure at 1600 C  for 0.5 h. Whereas, ZrB2 -10 vol% SiC and ZrB2-30 vol% SiC composite shows 65% more strength after exposure to 1500 C for 10 h, as reported by Guo and Zhang.11 By these reports, the improvement in strength is attributed to surface flaw healing by glassy phase. Our results show that the thickening of oxide scale has no influence on the strength of ZrB2 -SiC composites after oxidation at 1500 C for 5 h.        4.2.3. Exposure to 1700     C for 5 h                                The severity of oxidation of ZrB2-SiC composites for all contents of SiC at 1700 C is much more pronounced than 1500 C, as oxide scale  thickness after oxidation at 1700 C for 5 h was nearly 4-5  times more  than oxidation at 1500 C  for 5 h. The corresponding weight gain after oxidation at 1700 C  for 5 h was only 1.3-1.8  times higher  than after oxidation at 1500 C for 5 h. This happened due to loss of SiO2 in the form of gaseous SiO. Loss of SiO2 in form of SiO (g) is known as active oxidation. Active oxidation gets activated when the SiO2 layer grows, the partial pressure of oxygen reduces with oxide layer thickness and the SiO2 layer becomes unstable which undergo incongruent evaporation by formation of SiO (g) and O2 (g).25 The escape of SiO (g) along with B2O3 and CO gases from the samples leads to bubbles on the surface of ZrB2-SiC composites and oriented ZrO2 particles after exposure  to 1700 C for 5 h. Formation of bubbles at 1700 C shows  that  the vapor pressure of gaseous products is very high at 1700 C than 1500 C, as there were no bubbles observed on surfaces after exposure to 1500 C for 5 h. The orientation of ZrO2 particles was more pronounced in Z10S than Z30S since the glassy phase fraction was more in Z30S than Z10S. The glassy phase, which  is  liquid at 1700 C, separates the ZrO2 particles and prevents growth of ZrO2 particles.7 This results  into  less orientation of ZrO2 particles with higher SiC contents. A  dense  glassy SiO2 oxide  layer  formed  on  surface  of specimens after exposure 1700 C for 5 h (Fig. 14). The  thickness of dense glassy SiO2 oxide  layer was  found  to be more for ZrB2 -SiC composites with higher SiC contents. On other hand,  the  total oxide scale  thickness (thickness of SiO2 oxide layer + ZrO2 oxide  layer) was  thinner for higher SiC contents. But the residual strength variation with SiC contents has opposite  trends, which concludes  that both  surface flaws healing and generation of new defects  in oxide scale  thickness are not able  to explain  the behavior shown  for  residual strength after exposure  to 1500 C and 1700 C  for 5 h. Here, we are purposing a different mechanism, which explains  the behavior of residual  strength of ZrB2 -SiC composites with SiC contents after exposure 1700 C for 5 h. For  thicker oxide scale, cracks can generate  in oxide scale during bend  test. As  these cracks propagate  toward  the un-reacted ZrB2-SiC composites, cracks encounter the boundary between the oxide scale and un-reacted ZrB2-SiC composites. The interaction of cracks with the boundary depends upon the microstructure near the boundary. For the case of 1700 C  for 5 h, Z10S was  fully  sintered ZrO2 layer (Fig. 15a) whereas ZrO2 grain was separated by glassy phase                       \\x0c', '4466   M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467  Fig. 17. Fracture surface of ZrB2 -SiC composites after exposure to 1700     C for 5 h, (a) Z10S, (b) Z20S and (c) Z30S.  in Z20S and Z30S  (Fig. 15b) near  the boundary between  the oxide scale and un-reacted ZrB2 -SiC composites. Fully sintered ZrO2 layer can blunt the crack whereas the oriented ZrO2 grain, separated by glassy phase can propel  the crack  into un-reacted ZrB2-SiC composites. Blunting of crack at sintered ZrO2 layer, fracture surface of Z10S shows both intergranular and transgranular type fracture similar to as hot pressed condition (Fig. 17a). On other hand, only  intergranular fracture mode was observed on fracture surface of Z20S and Z30S (Fig. 17b and c). This is attributed in residual strength measurement, as very low residual strength for Z20S and Z30S composite.  5. Conclusions        ZrB2-SiC composites with 10-30 vol% SiC was hot pressed to  their  full density. The flexural strength of Z30S was  lower than Z20S. Oxidation  resistance was  increased with SiC content but  there was active oxidation of ZrB2-SiC composites at 1700 C. The effect of oxidation on residual strength was found to completely different for Z10S as compared to Z20S and Z30S. Due  to  lowest oxide scale  thickness of Z20S after exposure  to 1000 C for 5 hrs, surface flaws healing are supposed  to maximum. This  resulted  into maximum  improvement  in  residual strength. Surface healing effect for thicker oxide scale, formed after exposure  to 1500 C  for 5 h, was not evident  in  residual strength of ZrB2-SiC composites. After exposure  to 1700 C for 5 h, drastic degradation of residual strength was found for Z20S and Z30S composites. Whereas, Z10S has retained  their original strength after exposure  to 1700 C  for 5 h. A mechanism  is purposed for above behavior of residual strength after exposure in air to 1700 C, 5 h. The boundary between ZrO2 rich oxide scale and un-reacted ZrB2 -SiC composites has great influence on residual strength ZrB2-SiC composites at 1700 C. The formation of fully sintered ZrO2 layer at the boundary resulted into no degradation of residual strength foe Z10S.                 Acknowledgements  The authors  thankfully acknowledge  the financial  support received from DRDO, Govt. of India  in order  to carry out  the present  research study. We are also  thankful  to MBG, Tribology, NDTG EMG, and SFAG group of DMRL for bend  tests, hardness, Elastic modulus, SEM, XRD and EPMA experiments. We are grateful  to Director, Defence Metallurgical Research  Laboratory, Hyderabad  for giving permission  article and for his continuous support.  to publish   this  References  1. Upadhya K, Yang JM, Hoffmann WP. Materials for ultrahigh  structural applications. Am Ceram Soc Bull 1997;76:51-6.  temperature  2. Fahrenholtz WG, Hilmas GE. Talmy IG Zaykoski JA, Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90:1347-64.  3. Guo S. Densification of ZrB2 -based composites and  their mechanical and physical properties: a review. J Eur Ceram Soc 2009;29:995-1011.  4. Opeka MM, Talmy IG, Zaykoski JA. Oxidation-based materials selection for  2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci 2004;39:5887-94.  5. Karlsdottir SN. PhD thesis. University of Michigan, USA; 2007.  6. Eakins E, Jayseelan DD, Lee WE. Toward oxidation-resistant ZrB2 -SiC ultra  high  temperature  2011;42A:  ceramics. Metall Mater   Trans   A   878-87.  7. Zhang XH, Hu P, Han  JC. Structure evolution of ZrB2 -SiC during  oxidation in air. J Mater Res 2008;23:1961-2019.  the  8. Fahrenholtz WG. Thermodynamic analysis of ZrB2 -SiC oxidation. Formation of a SiC - depleted region. J Am Ceram Soc 2007;90:143-8.     9. Zhang H, Yan Y, Huang Z, Liu X,   effect of B4C   Jiang D.  Pressureless  sintering content. Scr Mater 2009;60:  of ZrB2 -SiC  559-62.  ceramics:   the   10. Guo SQ, Yang JM, Tanaka H, Kagawa Y. Effect of   thermal exposure on  strength of ZrB2 -based composites with nano-sized SiC particles. Compos Sci Technol 2008;68:3033-40.  11. Guo WM, Zhang GJ. Oxidation  resistance  and  strength  ZrB2 -SiC ceramics. J Eur Ceram Soc 2010;30:2387-423. 12. Zhi W, Qiang Q,  Zhanjan W, Guodeng  S.  Effect  of   retention   of  oxidation   at  1100 C on  the strength of ZrB2 -SiC-graphite ceramics. J Alloys Compd 2011;509:6871-968.     13. Patel M, Reddy JJ, Bhanu Prasad VV, Subrahmanyam J, Jayaram V. Residual  strength of hot pressed zirconium diboride  (ZrB2 ) after exposure  temperatures. Mater Sci Eng A 2012;535:189-96.  to high  14. Toraya H, Yoshimura M, Somiya S. Calibration curve for quantitative anal ysis of  the monoclinic-tetragonal ZrO2 Ceram Soc 1984;67:C119-21.  system by X-ray diffraction. Am  15. Anstis GR, Chantikul P, Lawn BR, Marshal DB. A critical evaluation of  indentation  techniques  for measuring  fracture  toughness:  measurements. J Am Ceram Soc 1981;64:553-638.  I, direct crack  16. Zhang SC, Hilmas GE, Fahrenholtz WG. Mechanical properties of sintered ZrB2 -SiC ceramics. J Eur Ceram Soc 2011;31:893-901. 17. Chamberlain AL,  Fahrenholtz WG, Hilmas GE,  Ellerby DT. High strength zirconium diboride-based ceramics. J Am Ceram Soc 2004;87:  1170-2.  18. Mallik M, Ray KK, Mitra R. Oxidation behavior of hot pressed ZrB2 -SiC and HfB2 -SiC composites. J Eur Ceram Soc 2011;31:199-215. 19. Tripp WC, Davis HH, Graham HC. Effect of an SiC addition on the oxidation  of ZrB2 . Am Ceram Soc Bull 1973;52:612-3.  \\x0c', '20. Rezaie A, Fahrenholtz WG, Hilmas GE. Evolution of structure during   the  23. Karlsdottir SN, Halloran   oxidation of zirconium diboride-silicon carbide in air up to 1500 Ceram Soc 2007;27:2495-501.  C. J Eur  content  on  solid  2009;92:481-6.  and      M. Patel et al. / Journal of the European Ceramic Society 32 (2012) 4455-4467   JW. Oxidation of ZrB2 -SiC:  liquid  oxide  phase  formation.   4467  influence of SiC  J Am Ceram   Soc  21. Gangireddy S, Karlsdottir SN, Halloran JW. Liquid oxide flow during oxida 24. Gulbransen EA, Jansson SA. The high   temperature oxidation,   reduction,  tion of zirconium diboride-silicon carbide ultra high temperature ceramics. Key Eng Mater 2010;434-435:144-8.  and volatilization reactions of silicon and silicon carbide. Oxidation of metal 1972;4:181-91.  22. Karlsdottir SN, Halloran JW, Grundy AN. Zirconia transport by liquid con 25. Heuer AH, Lou VLK. Volatility diagrams   for   silica,   silicon nitride, and  vection during oxidation of zirconium diboride-silicon carbide. J Am Ceram Soc 2008;91:272-7.  silicon carbide and their application to high temperature decomposition and oxidation. J Am Ceram Soc 1990;73:2785-828.  \\x0c']"
},{
  "_id": 248,
  "PDF": "Strong ZrB2-SiC-WC Ceramics at 1600°C.pdf",
  "Text": "['-SiC-WC Ceramics at 1600°C Strong ZrB2  Ji Zou,  ‡,§,¶  Guo-Jun Zhang,  ‡,†  Chun-Feng Hu,  -  Toshiyuki Nishimura,  ¶  Yoshio Sakka,  ¶,†  Jef Vleugels,  §,†  and Omer Van der Biest  §  ‡  State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics,  Shanghai 200050, China  §  Department of Metallurgy and Materials Engineering (MTM), Katholieke Universiteit Leuven, B-3001 Heverlee, Belgium  ¶  National Institute for Materials Science NIMS, 1-2-1 Sengen, Tsukuba Ibaraki, 305-0047, Japan  High  ZrB2-20 vol% SiC temperature flexural strength of (ZS) up to 1600°C in high purity argon atmosphere  ceramics  was  significantly  improved  by  adding 5 vol% WC, but added. ZrB2-20SiC-5WC ceramic (ZSW) has a very high strength (mean ± SD) of 675 ± 33 MPa at 1600°C, and also an elastic and transgranu degraded when  5 vol% ZrC was  lar  fracture mode was observed. According to the analysis of  the  fracture modes  and  crack origins in ZSW ceramics, 1000°C was attributed to  the  improvement  in  strength above  the  removal of the oxide impurities from grain boundaries.  I.  Introduction  Z IRCONIUM  diboride,ZrB2, forced by various amounts of SiC (ZS) possess many advantages1-3: high eutectic (~2310°C), relative (500-1000 MPa)4, high room temperature flexural (17-20 GPa), hardness 1500°C,5,6 tance below \\x001]7 [60-100 W·(m·K) as well. The these properties makes ZS as a promising candidate for aero particulate  composites  rein temperature  strength  and  high  excellent  oxidation  resis and  high  thermal  conductivity  unique  combination  of  space applications such as  leading edges and thermal protec tion systems for reusable atmospheric hypersonic flight vehicles.1-3 In these applications, ZS ceramics need to have  re-entry  vehicles and  thermal,  mechanical, and chemical stability at ultra high temperatures (>1600°C) and under dissociated oxygen. Therefore, enhancing the high temperais a key for utilizing ZrB2-SiC composites practice. Despite numerous investigations regarding process low partial pressure of molecular, and  ture properties  in  ing and high temperature properties of ZS, e.g., oxidation/ ablation resistance,5,6 thermo-physical7 and creep property,8  were  reported in the  last decades,  the  research activity con cerning the high temperature flexure lacking.9,10 The highly covalent nature of B-B, Si-C bond and partial covalent bonding of Zr-B result in high bending strength of ZrB2-SiC ceramics at room temperature.2 However, as with most of the engineering ceramics,  strength of ZS is  still  the flexure strength of  the  ZS  ceramics dropped down was over 1000°C-1300°C as Loehman11 and Rhodes12  rapidly when  test temperature literature.11,12  reported  in  the  reported that  the flexural  strength  of ZS at 1400°C only maintained 48% and 83% of sured at room temperature, respectively. As we know, oxide  that mea impurities  either  from raw materials or  the processing tend  to concentrate at  the grain boundaries (wetting grains) or tri ple junction points  (non-wetting grains) during densification.  The  existence of  these  low melting point  secondary phases  will adversely affect  the grain boundary strength especially at  high temperatures.  The purpose of  this article  is  to explore  the  feasibility to  minimize  (or  even  eliminate)  the  low melting  point  oxide  phases  at  the  grain boundaries of ZS ceramics using high  purity powders and a proper  selection of  sinter additives  in  view to maximize their high temperature strength.  II.  Experimental Procedure  The  characteristics of  the  raw powders  and their  suppliers  are  summarized in Table I. High purity ZrB2 powder was synthesized through a boron/carbon thermal reaction  between ZrO2 and B4C in vacuum (similar to HfB2 powder in Ni’s work, i.e., sample HB2 synthesized at 1600°C for 1.5 h in vacuum13), other powders  are  commercially  avail able. ZrB2 5 vol% MC powders  and  SiC (20 vol%) powders with and without (M = W and Zr) were ball-mixed in  acetone for 12 h in a resin jar using Si3N4 mixing media. The samples of ZrB2-SiC without and with WC or ZrC addition were marked as ZS, ZSW, and ZSZ, respectively. After mix ing,  the  slurry  of  the  powder mixture was  dried  through  rotary evaporation and sieved to 200 mesh. The compacts (6 mm 9 30 mm 9 37 mm, height 9 width 9 length) of the mixed powders were hot pressed at 1900°C (for ZS at 2000°C)  and 30 MPa for 1 h. The furnace atmosphere consisted of a mild vacuum (~20 Pa) for temperatures up to 1650°C. Above 1650°C the atmosphere was details of the densification process were similar to PA program as reported in our previous works.14  switched to flowing argon. The  The  test  bars with  dimensions  of  2 mm 9 2.5 mm by  25 mm were cut from the sintered pellet. One tensile surface of each specimen was polished and finished down to 1.5 lm  whereas  the  edges were  chamfered. Flexure  strengths were  measured using a three-point bending configuration ranging to 1600°C at a cross-head speed of 0.5 mm/min (Model 4505; Instron Corp., Norwood, MA).  temperature  from ambient  To limit  the effect of  the oxidation during measurements, \\x003 Pa, evacuated down to 10 argon at a flowing  the  testing  chamber was first  and  then it was backfilled with high purity  rate  of  100 mL/min.  The temperature was raised to the rate of 30°C/min. To reach a thermal  desired temperature at  equilibrium,  the  load was  applied  after  the  specimen was  soaked at  the testing temperatures  for 20 min. The reported  F. Monteverde—contributing editor  Manuscript No. 30577. Received October 30, 2011; approved December 16, 2011.  †  Authors  to whom correspondence should be addressed. e-mails: gjzhang@mail.sic.  ac.cn, sakka.yoshio@nims.go.jp, and jozef.vleugels@mtm.kuleuven.be  874  J. Am. Ceram. Soc., 95 [3] 874-878 (2012)  DOI: 10.1111/j.1551-2916.2011.05062.x  © 2012 The American Ceramic Society  Journal  \\x0c', 'flexural  strengths are the average of five specimens  for  room  temperature  and of  three  specimens  for high temperatures.  The microstructures were  characterized using X-ray diffrac tion (XRD; D/max 2550V, Rigaku, Tokyo, Japan) and scan ning  electron microscopy  (SEM; TM3000, Hitachi, Tokyo,  Japan) before and after bending test.  III.  Results and Discussions  The measured bulk densities of ZS, ZSW, and ZSZ are 5.53, and 5.60 g/cm3,  6.01,  respectively. Relative densities of  the  three ceramics are higher than 99% and no apparent residual  porosity was  found on their polished surfaces observed by  SEM (Fig. 1). Porosity is  thus not expected being the origin  for  the  strength  differences  at  room temperature  for  these  samples. From Fig. 1(a), it can be found that some platelike SiC grains with 5.3 ± 2.1 lm in length and 1.1 ± 0.8 lm in  width were homogeneously dispersed in ZrB2 matrix. However, the morphology of SiC grains in ZSW and ZSZ was quasi-spherical with average particle size of 1.1lm [Figs. 1(b)  and (c)]. Some  secondary phases  in ZSW were  identified as  (Zr, W)ssC [A arrowed arrowed in Fig. 1(b)] respectively, according electronic diffraction in our previous work.15,16 The room-temperature flexural strength (mean ± SD) varied from 546 ± 55 MPa for ZS to 1101 ± 127 MPa for ZSZ.  in  Fig. 1(b)]  and  (W,  Zr)ssB to XRD and  [B  Previous  investigations  showed that  the average  size of  the  SiC  grains  was an important factor in strength of the ZrB2-SiC ceramics. The room temperature strength of ZS increased as SiC particle size decreased.17 So  determining  the  the  lowest  room-temperature  strength of ZS was attributed  to its  largest SiC grain size. Flexural  strength of ZS, ZSW, 1600°C is  and ZSZ as  a  function  of  temperatures  up  to  shown in Fig. 2(a). The ZSZ material has average room tem perature strength of over 1 GPa. However, its strength drastically decreased when temperature was over 1000°C and the  strength of ZSZ further decreased to the  lowest one,  com pared to that of ZS and ZSW, when temperature was over 1300°C. The strength of ZS slightly increased with tempera1000°C  ture  increasing  to  a maximum strength  at  around  presumably due  to the  crack healing on the  tensile  surfaces  of  bars  during  the  high  temperature  test,  and  then  it  decreased  with  temperature.  ZSW exhibited 1300°C, strength at temperatures above and no further strength degradation up to 1600°C was observed. The flexural strength (mean ± SD) of ZSZ, ZS, and ZSW measured at 1600°C are 317 ± 7, 460 ± 31, and 675 ± 33 MPa, which  the  highest  are  28.7%,  84.2%,  and 111.5% of  their  room-temperature  strength, respectively.  Considering that starting ZrB2 and SiC powders, equipment and processing procedures are similar  sintering  for ZSZ  and ZSW,  the differences in the high temperature strength of  these samples should mainly depend on the carbide additives  added in the different  samples. This  is because  the  carbide  additions will determine whether  the grain boundary phases  exist or not and the  characters of  the grain boundaries.  In  addition, the carbides inhibited the grain growth of -SiC matrix during the sintering process.  the ZrB2  The average of both ZrB2 and SiC grain sizes in ZSZ and ZSW are similar, so the difference in strength at high temper Table I.  Characteristics of the Raw Materials  Powders  Particle size (D50), lm  Impurities (wt%)  Supplier  ZrB2  1.05  O 0.46, C 0.10, Hf 1.10, Fe 0.08, Ca 0.06, Ti 0.016, Others <0.01  Home-made  a-SiC  0.45  B 0.33, O 1.00, Ca 0.24, Cl 0.10,  Fe 0.16, V 0.09  Changle Xinyuan Carborundum Co. Ltd  ZrC  ~2  Hf 3.0, Ti 0.13, Nb 0.05, O 0.78, Mg 0.10  High Purity Chemical Institute Co. Ltd.,  Saitama, Japan  WC  ~1  C 6.13, Cr 0.03, Co 0.01, Mo 0.01, O 0.20  Hard alloy Co. Ltd, Zhuzhou, China  (a)  (b)  (c)  (d)  (e)  (f)  Fig. 1. Microstructures of sintered ZS (a), ZSW (b), and ZSZ (c). Some typical micrographs on the tensile surfaces of bars (d) ZS, (e) ZSW (f) ZSZ near the fractured parts after high temperature bending tests at 1600°C. The black, gray, and white phases in (c) are SiC, ZrB2, and ZrC, respectively.  and,  March 2012  Rapid Communications of  the American Ceramic Society  875  \\x0c', 'ature  should be  caused by their different boundary phases.  The morphology of ZSZ and ZSW fracture surfaces tested at 1600°C is shown in Fig. 3. The surface macro-topography of ZSZ is very smooth, so that it is hard to distinguish the ten sion side and the compression side [Fig. 3(g)]. At  room tem perature, both ZSZ and ZSW showed fracture behavior. At 1600°C,  a  transgranular  the fracture mode of  the ZSZ  material  [Figs. 3(h) and (i)] changed to a fully intergranular  fracture mode.  In  contrast,  the  fracture mode  of  ZSW  remained transgranular [Figs. 3(e) and (f)]. Moreover, the surfaces of ZSW fractured at 1600°C have a radiating pat tern that indicates the crack propagation direction [Fig. 3(d)].  Such observations demonstrate that the fracture behavior of ZSW at 1600°C remains brittle, whereas in the ZSZ material partial plastic deformation took place. This conclusion could  also be deduced from the load-displacement curves as shown  in Fig. 2(b). The load-displacement plot  for ZSW fits nicely;  a linear behavior before the fracture point,  i.e.,  the Young’s  modulus of ZSW almost kept constant during the testing at 1600°C. However,  the  load-displacement plot  for ZSZ devi ated from linearity and its Young’s modulus was gradually  reduced with  time. The  phenomena  confirmed  again  that  ZSW showed a brittle fracture behavior whereas ZSZ plastically deformed during the testing at 1600°C.  The  fracture of bars during bending test  is  the  result of  competition between transgranular strength and intergranular  strength  in  a  certain  sample. High  transgranular  strength  leads  to inter-granular  fracture, while  transgranular  fracture  is  predominant when  the  intergranular  strength  is  higher.  The  intra-granular  strength of ZrB2 posed to be the same with ZrB2 grains  grains  in ZSZ is  sup in ZSW as they have  similar average grain size. According to the fracture modes, the sequence in strength at 1600°C should be: transgranular strength of ZSW > intra-granular strength of ZSW > ZSZ intergranular strength transgranular strength of ZSZ. A  >  clear  indication for  this  sequence is  that  the grain boundary  of ZSZ was easier to be etched, compared to that of ZSW, after the same heat treatment during the bending test [1600°C,  20 min, Figs. 1(e) and (f)].  Why do WC and ZrC additions have opposite  roles on  the  high  temperature  strength  of  ZS matrices?  It  was  reported that  the bending strength of  the ZrB2-based composite at elevated temperatures mainly depends on the glassy  oxide  content  and grain size of  the  composites, because  a  high content of glassy oxide and a small particle size lead to grain boundary sliding.10,18 For the current ZrB2-SiC system, oxide impurities are present on the surfaces of the starting  powders which include ZrO2, amorphous B2O3, and SiO2. Among these oxides, only B2O3 can be easily removed by evaporation at elevated temperatures under  19  vacuum during (~15 Pa  the sintering process due to its high vapor pressure at 1300°C). can react with ZrO2 through reaction (1) and (2) removing it (<1650°C).16,20 WC can also react with at mild temperatures  In previous works,  it has been proved that WC  SiO2 culations, the favorite is 1115°C, which is lower than the isothermal temperature (1650°C) during the sintering process. The above discussions indicate that WC can remove the oxide impurities in ZrB2-SiC system and clean grain boundaries in ZSW are therefore  through reaction (3). According to thermodynamic cal temperature  for  reaction (3) at 10 Pa  expected:  this was already shown using high-resolution transZrB2-SiC-WC  mission electron ceramics.16  microscopy  analysis  in  3WC þ ZrO2 ¼ ZrC þ 3W þ 2COðgÞ  ð1Þ  3ZrB2 þ 6WC þ ZrO2 ¼ 4ZrC þ 6WB þ 2COðgÞ  ð2Þ  WC þ SiO2 ¼ W þ SiOðgÞ þ COðgÞ  ð3Þ  Similar thermodynamic calculations were performed in the  case of ZrC used as sinter additive. However, none of favorable below 2000°C in vacuum,  these  reactions  is  i.e., oxide  impurities  could  not  be  removed  during  densification:  this  should be one reason for the strength degradation of ZS and  ZSZ at elevated temperatures.  In addition,  in the  fractured parts of ZS and ZSZ,  it  is  easy to find some  cavitations  located at  the  triple points of  SiC and ZrB2 grains these cavities should be associated with the grain sliding durthe bar, as verified by Guo et al.21 Grain  [Figs. 1(d) and (f)]. The  formation of  ing the loading of  sliding occurs when the stresses could not be accommodated  by  homogeneous  plastic  deformation  of  grains. Thus  SiC  grain has to rotate combining with cracks nucleated ahead of  their  edges, a result of  such coalescence of  cracks was new  cavities formed. Strong grain boundaries decrease the rate of  grain sliding,  so no cavities were  found in ZSW [Fig. 1(e)].  Once  cavities were  initiated,  the  stresses  that  concentrated  around the edges of  the cavities will make the cracks to grow  and propagate perpendicularly to the loading direction, caus ing the specimens  to exhibit  intergranular  fracture as  shown  in Figs. 3(b) and (h). The presence of cavities is another rea son why the strength of the ZS and ZSZ materials decreased drastically above 1000°C.  Based  on  the  load-displacement  curve  [Fig. 2(b)],  the  decrease in fracture strength for ZSZ was mainly controlled  by its remarkable plastic deformation at higher temperatures.  In fact,  apart  from the  smaller  grain size  in ZSZ,  another  factor  that  should account  for  the  strength drop is  that  the  Peierls  shear  stress  to initiate the movement of a dislocation  (a)  (b)  Fig. 2.  (a)  The  flexural  strength  of  ZS, ZSW, and ZSZ load-displacement  as  a  function of testing temperature, for ZS, ZSW, and ZSZ at 1600°C.  (b) The  curves  876  Rapid Communications of  the American Ceramic Society  Vol. 95, No. 3  \\x0c', 'March 2012  Rapid Communications of  the American Ceramic Society  877  (a)  (d)  (g)  (b)  (e)  (h)  (c)  (f)  (i)  Fig. 3.  The flexural surfaces of ZS, ZSW, and ZSZ at 1600°C.  on glide plane of cubic ZrC is obviously lower  than that of  hexagonal ZrB2.  22  IV.  Conclusions  The present work showed that a ZrB2-SiC based UHTC had no strength degradation up to 1600°C in inert atmosphere by  using high purity raw powders and adding small amounts of  WC as sintering aid. The average flexural strength tested at 1600°C was as high as 675 MPa, which is 111.5% of  its flex ural strength at room temperature. Grain sliding, which is the main reason for the strength decrease of ZrB2-SiC based composites above 1000°C, can be effectively limited by  designing  larger  average  grain size  and/or  glassy  free  grain  boundaries. Current  result provides  a new way  to improve  the strength of UHTCs at elevated temperatures.  Acknowledgments  This work was financially supported by NSFC (No. 50632070 and 91026008),  the bilateral project of NSFC-JSPS (No. 51111140017),  the Research Fund of  K.U. Leuven under project GOA/08/007. The Chinese Academy of Sciences  Hundred Talents Program is gratefully acknowledged. Ji Zou Research Fund of K.U. Leuven for his post-doctoral F+/11/006  thanks  the  fellowship.  The authors also thank Prof. S. C. Zhang and the reviewer for their great help  in the discussion and revision of  the draft.  References  1K. Upadhya, J.-M. Yang, and W. P. Hoffman, “Materials Temperature Structural Applications,” Am. Ceram. Soc. Bull.,  for Ultrahigh 51-6  [12]  76  (1997). 2W. G. Fahrenholtz and G. E. Hilmas, “Refractory Diborides of Zirconium and Hafnium,” J. Am. Ceram. Soc., 90 [5] 1347-64 (2007).  3M. Opeka,  I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey,  “Mechanical, Thermal and Oxidation Properties of Refractory Hafnium and Zirconium Compounds,” J. Eur. Ceram. Soc., 19 [13-14] 2405-14 (1999). 4A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby, “High-Strength Zirconium Diboride-Based Ceramics,” J. Am. Ceram. Soc., 87 [6] 1170-2 (2004). 5A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Oxidation of Zirconium Diboride-Silicon Carbide at 1500°C at a Low Partial Pressure of Oxygen,” J. Am. Ceram. Soc., 89 [10] 3240-5 (2006). 6F. Monteverde and R. Savino, “Stability of Ultra-High-Temperature ZrB2- SiC Ceramics Under Simulated Atmospheric Re-Entry Conditions,” J. Eur. Ceram. Soc., 27 [16] 4797-805 (2007). 7J. W. Zimmermann, G. E. Hilmas, and W. G. Fahrenholtz, “Thermophysical Properties of ZrB2 and ZrB2-SiC Ceramics,” J. Am. Ceram. Soc., 91 [5] 1405-11 (2008). 8I. G. Talmy, J. A. Zaykoski, and C. A. Martin, “Flexural Creep Deformain Oxidizing Atmosphere,” J. Am. Ceram. Soc., 91  tion of ZrB2/SiC Ceramics [5] 1441-7 (2008). 9D. Kalish, E. V. Clougherty, and K. Kreder, “Strength, Fracture Mode, and Thermal Stress Resistance of HfB2 and ZrB2,” J. Am. Ceram. Soc., 52 [1] 30-6 (1969). 10P. Hu and Z. Wang, “Flexural Strength and Fracture Behavior of ZrB2- SiC Ultra-High Temperature Ceramic Composites at 1800°C,” J. Eur. Ceram. Soc., 30, 1021-6 (2010). 11R. Loehman, E. Corral, H. P. Dumm, P. Kotula, and R. Tandon, “Ultra Sandia  High Temperature Ceramics Report Sand., 2006-925 (2006). 12M. J. Gasch, D. T. Ellerby, and S. M. Johnson, “Ultra High Temperature Ceramic Composites”; pp. 212-4 in Hand Book of Ceramic Composites, Edited  for Hypersonic Vehicle Applications,”  by N. P. Bansal. Kluwer Academic Publishers, Norwell, MA, 2005. 13N. D. Wei, G. J. Zhang, Y. M. Kan, and P. L. Wang, “Synthesis of Mon odispersed Fine Hafnium Diboride Powders Using Carbo/Borothermal Reduction of Hafnium Dioxide,” J. Am. Ceram. Soc., 91 [8] 2709-12 (2008). 14J. Zou, G.-J. Zhang, Y.-M. Kan, and P.-L. Wang, “Hot-Pressed ZrB2-SiC Ceramics with VC Addition: Chemical Reactions, Microstructures and Mechanical Properties,” J. Am. Ceram. Soc., 92 [12] 2838-46 (2009). 15J. Zou, G. J. Zhang, and Y. M. Kan, “Formation of Tough Interlocking  Microstructure in ZrB2-SiC Based Ultra High Temperature Ceramics by Pressureless Sintering,” J. Mater. Res., 24 [7] 2428-34 (2009). 16J. Zou, S.-K. Sun, G.-J. Zhang, Y.-M. Kan, P.-L. Wang,  and T. Ohji,  “Chemical Reactions, Anisotropic Grain Growth and Sintering Mechanisms  \\x0c', '878  Rapid Communications of  the American Ceramic Society  Vol. 95, No. 3  of Self-Reinforced ZrB2-SiC Doped with WC,” 1575-83 (2011). 17S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, “Mechanical ProperSintered ZrB2-SiC Ceramics,” 893-901 J. Eur. Ceram. Soc.,  J. Am. Ceram. Soc.,  ties  94  [5]  of  31  [5]  (2011). 18A. Bellosi, F. Monteverde, and D. Sciti, “Fast Densification of Ultra-High Temperature Ceramics by Spark Plasma Sintering,” Int. J. Appl. Ceram. Technol., 3 [1] 32-40 (2006). 19S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, “Pressureless Sintering of ZrB2-SiC Ceramics,” J. Am. Ceram. Soc., 91 [1] 26-32 (2008).  20S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, “Zirconium Carbide- In Situ Reaction Sintering,” J. Am. Ceram.  Tungsten Cermets Prepared by Soc., 90 [6] 1930-3 (2007). 21W. M. Guo, G.-J. Zhang, and H.-T. Lin, “High-Temperature Flexural in Argon Atmosphere,” Ceram Int., 38 [1] 831-5  Creep of ZrB2-SiC Ceramics  (2012). 22J. S. Haggerty and D. W. Lee, “Plastic Deformation of ZrB2 Single Crystals,” J. Am. Ceram. Soc., 54 [11] 572-6 (1971).  h  \\x0c', \"Copyright of Journal of the American Ceramic Society is the property of Wiley-Blackwell and its content may  not be copied or emailed to multiple sites or posted to a listserv without the copyright holder's express written  permission. However, users may print, download, or email articles for individual use.  \\x0c\"]"
},{
  "_id": 249,
  "PDF": "Structural and compositional analyses of oxidised layers of ZrB2-based UHTCs.pdf",
  "Text": "['Journal of the European Ceramic Society 35 (2015) 4059-4071  Contents lists available at www.sciencedirect.com  Journal of the European Ceramic Society  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j e u r c e r a m s o c  Structural and compositional analyses of oxidised layers of ZrB2-based UHTCs  D.D. Jayaseelan ∗ , E. Zapata-Solvas, R.J. Chater, W.E. Lee  Centre for Advanced Structural Ceramics, Department of Materials, Imperial College, London SW7 2AZ, UK  a r t i c l e  i n f o  a b s t r a c t  ZrB2 based UHTCs (ZrB2 /20 vol.% SiC) with rare earth additions (La2O3 , LaB6 and Gd2O3 ) were oxidised for 1 h at 1600   C and the oxidised layers analysed by FIB-SIMS and TEM techniques. Rare earth elements added to ZrB2 /20 vol.% SiC led on oxidation to formation of a thick (up to 350 \\u242em) protective oxide layer. Detailed cross-sectional analyses of oxidised layers revealed that they comprised of mixed crystalline and amorphous phases: m-ZrO2 and zirconates were formed on the oxidised surfaces in all samples as well as evidence of immiscible liquid.  © 2015 Elsevier Ltd. All rights reserved.  Article history:  Received 26 June 2015 Received in revised form 17 July 2015 Accepted 20 July 2015 Available online 15 August 2015  Keywords:  UHTC Oxidation FIB-SIMS TEM Characterisation  1.  Introduction  The development of oxidation resistant ultra high temperature ceramics (UHTCs) is being increasingly studied because of a growing need for materials to operate in extreme environments in aerospace applications [1-14]. ZrB2 and HfB2 ceramics are candidate materials for UHTCs applications owing to their unique combination of properties, including high melting point (>3000  C), high strength, high elastic modulus and high thermal conductivity. Because monolithic borides are vulnerable to oxidation attack, the state-of-art UHTCs are ZrB2 /SiC [13,15-22] and HfB2 /SiC [1,23-25] composites. Addition of SiC to ZrB2 or HfB2 -based ceramics not only improves their mechanical properties but also their oxidation resistance by altering the properties of the glassy phase in the oxidised layer. High levels of silica in the borosilicate glassy phase increase its viscosity and therefore reduce the diffusion of oxygen through the liquid. Systems with higher viscosity and increased liquidus temperatures have the added benefit of suppressed evaporation of boria from the glassy phase. Diffusivity is inversely proportional to the viscosity of the liquid through which the diffusion is taking place, shown in the Stokes-Einstein relationship [26]:  D =  kT  6(cid:2)\\x01r  ∗ Corresponding author. Tel.: +44 2075895111x51181. E-mail address: d.j.daniel@imperial.ac.uk (D.D. Jayaseelan).  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.07.026 0955-2219/© 2015 Elsevier Ltd. All rights reserved.  [m2 s−1 ], where D = diffusion constant k = Boltzmann’s constant [J K−1 ], T = temperature [K], \\x01 = viscosity [m2 s−1 ], and r = spherical particle radius [m]. The viscosity of the borosilicate glass can be further increased by addition of certain elements to the bulk material. Talmy et al. [27] found that the oxidation resistance of hot pressed ZrB2 + 25 vol% SiC composites could be improved by addition of diborides of Cr, Ti, Ta, Nb and V. These additions result in production of the respective oxides in the glass. The improvement in oxidation resistance comes from the fact that borate and silicate glasses containing oxides of the Group IV-VI transition metals are immiscible and therefore lead to phase separation. Such systems contain compositions with very high viscosity and liquidus temperatures [26]. The immiscibility of a glass increases with increasing cation field strength of the metallic oxide-forming element, z/r2 , where z is the valence and r is the ionic radius. The order in which the oxidation resistance was improved correlated well with the cation field strength of the modifying additive. Further work by Talmy et al. [28] took a novel approach by producing ceramics in the systems ZrB2 -Ta5 Si3 . Ta5 Si3 provides the tantalum to induce glass immiscibility and silicon to form the protective borosilicate glass in the oxidised layer. In the same way, addition of 10 and 20 vol.% tungsten carbide to ZrB2 -SiC ceramics increased the oxidation resistance of ZrB2 ceramics by changing the morphology of the zirconia scale formed during oxidations [29-31]. The physical and chemical processes that occur at the exposed surface depend upon the microstructure and composition of the oxidised material. It follows that the modification of microstructure and composition can have a beneficial (or detrimental) effect  \\x0c', '4060  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  on materials’ oxidation resistance. Previously we reported oxidation protective refractory coatings formed by in situ processing of rare earth additive containing 20 vol.% SiC reinforced ZrB2 during oxidation at 1600  C [13]. Subsequently, we studied the long-term oxidation kinetics of UHTCs with La2O3 added at 1400-1600  C and analysed the protectiveness of the oxide products formed during oxidation [32]. In the present study, the structure and composition of the oxide layers on UHTCs modified by the addition of rare earth elements were studied using nano analytical characterisation techniques and the reaction pathways are discussed.  2. Experimental details  ZrB2 (>99%, (cid:4) = 6.085 g/cm3 , Sigma Aldrich, Gillingham, UK) and SiC (␣-SiC, 99%, (cid:4) = 3.217 g/cm3 , Good Fellow Chemicals, Huntingdon, UK) were used as starting materials to form base line ZrB2 /20 vol.% SiC (hereafter termed ZS20) UHTC. Rare earth (RE) additions in the form of La2O3 (>99%, (cid:4) = 6.51 g/cm3 , Fluka, USA), LaB6 (>99%, (cid:4) = 4.72 g/cm3 , Fluka, USA) and Gd2O3 (>99%, (cid:4) = 7.41 g/cm3 , Fluka, USA) were added to the starting materials. UHTCs namely, ZS/10 wt.% La2O3 (ZSLO), ZS/10 wt.% LaB6 (ZSLB) and ZS/10 wt.% Gd2O3 (ZSGO) were densified in a spark plasma sintering (SPS) furnace (FCT Systems, Germany) using the sintering parameters given in Table 1 [13]. Oxidation studies were performed for the sintered (99% dense) samples in a bottom loaded open hearth MoSi2 furnace by heating samples to 1600  C at a rate of 10  C/min,  Table 1 Sintering parameters and density of all compositions.  Temp.  Heating rate  Pressure  Dwell time  ZSLO ZSLB ZSGO  1850   C 1750   C 1800   C  100   C/min 100   C/min 100   C/min  70 MPa 70 MPa 70 MPa  5 min 5 min 5 min  Relative density 99% 99% 99%  holding for 1 h and finally air-quenching to room temperature by lowering the specimen chamber. Oxidised samples were cut, mounted in epoxy resin moulds and polished down to 1 \\u242em surface finish. Microstructures of the sintered and oxidised samples were observed under a scanning electron microscope (SEM) fitted with a field-emission gun (FEG, model LEO15) and energy dispersive spectroscopic (EDS, X-Max 20, Oxford instruments; Abingdon, UK) elemental analysis was performed to assist phase identification. Additionally, oxidised samples were imaged and oxidised layers compositionally analysed in an FEI FIB200T-SIMS focussed ion beam workstation (FIB) with secondary ion mass spectroscopy facility attached (SIMS). Gallium ions were used to sputter the sample surface. The samples were attached to an aluminium sample holder with the help of silver tape to avoid charge build-up on the sample surface that would affect the secondary ion detection. TEM sections on the oxide layers were prepared using the FIB workstation operating with a gallium beam  Fig. 1. BF TEM images of ZS20 with RE additives (a) ZS20/10% La2 O3 , (b) ZS20/10% LaB6 and (c) ZS20/10% Gd2 O3 .  \\x0c', 'D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  4061  at 30 keV. Microstructural and chemical analyses were also carried out using a JEOL 2000FX operating at 200 kV with an Oxford Instruments microanalysis attached. Selected Area Electron Diffraction (SAED) patterns were indexed using single crystal diffraction patterns from Williams and Carter [9] and TEM diffraction analysis software (SingleCrystal, CrystalMaker Software Ltd.; Begbroke, UK)  grains. In addition to the ZrB2 , SiC, and Gd2O3 starting materials, other phases present are ZrO2 (darkest contrast) and gadolinium silicate (likely formed by reaction of silica on the SiC with the Gd2O3 on sintering and whose morphology suggests it is glass). ZrO2 could be the contamination from the grinding media or have arisen from the oxygen impurities associated with ZrB2 powders.  3. Results and discussion  3.1.  Sintering and microstructure  All samples, ZSLO sintered at 1850  C, ZSLB sintered at 1750  C, and ZSGO sintered at 1800  C, attained above 99% of theoretical density in less than 5 min. Phase analysis confirmed that all starting materials remained unchanged after SPS and did not show any secondary crystalline phase analysed by XRD [13]. Fig. 1 shows typical bright-field TEM images of (a) ZSLO, (b) ZSLB and (c) ZSGO. Fig. 1(a) shows the distribution of SiC and ZrB2 grains in ZSLO. No grain boundary phases were observed. Previous study suggested that during liquid phase sintering of SiC, La2O3 reacts with the surface SiO2 on SiC to form a lanthanum silicate phase [13]. Although we have no direct evidence for this in the areas analysed, it seems likely that at the high sintering temperature, 1800  C, used in the present study at least some of the RE oxide would have reacted with the surface SiO2 on SiC to form a liquid grain boundary silicate (RE2 Si2O7 ) phase. Fig. 1(b) shows the distribution of ZrB2 , SiC and LaB6 grains in ZSLB. Intergranular SiC grains were either 6H or 15R polytype and elongated (1-2 \\u242em length and 0.5 \\u242em diameter), whereas the intragranular SiC was cubic 3C polytype. ZrB2 grains have faceted morphology. Fig. 1(c) shows the bright-field TEM image of ZSGO whereby SiC grain is surrounded by Gd2O3  3.2. Oxidation  3.2.1. Oxidation of ZS20/10 wt.% La2O3 (ZSLO)  Fig. 2(a) is an electron image (EI) of cross-section surface microstructures of oxidised ZSLO at 1600  C taken in FIB station. Fig. 2(a) reveals three oxidised layers. The outer >250 \\u242em thick layer (1) is predominantly a continuous bright phase with isolated (5-30 \\u242em diameter) dark regions present. The intermediate layers above the unaffected bulk material are about 50 \\u242em thick. Two intermediate layers can be distinctly seen in this FIB electron image. SIMS data collected (not shown here) on the dark contrast phase (layer 2) in Fig. 2 (b) consists mainly of Si and O whereas the bright phase (layer 3) is predominantly Zr and O. SIMS analysis was carried out at the interface between the top surface layer and the top intermediate layer in Fig. 2(a). The crater formed after SIMS analysis is shown in Fig. 2(b) revealing clearly the presence of two different phases. The contrast of the ion signal can be different from the SE contrast (channelling, voltage contrast) and hence ion imaging can give additional information. SIMS analysis (Fig. 2(c)) revealed the presence of Si, Zr, La, B and their oxides in the oxidised top surface. Interestingly, boron was also detected indicating that in ZSLO, even after oxidation for 1 h at 1600  C, boron was not completely removed. Reports on the detection of boron in the oxidised layers using EDS in the SEM are sparse with limited ability to detect boron.  (a) FIB-EI of cross-section microstructure of ZSLO oxidised at 1600   C for 1 h showing 250 \\u242em thick oxidised layers, (b) crater formed at the interface of layers 1 and Fig. 2. 2 and (a) and (c) positive mass spectra taken on rectangular crater of (b) showing the presence of B, Si, Zr, La and their respective oxides.  \\x0c', '4062  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  Fig. 3. The electron image of an ion polished region of the oxidised layer in ZSLO obtained at a tilt angle of 45  craters showing the presence of B, Si, Zr, La and their respective oxides.  in the FIB workstation. SIMS analysis carried out at circular  Since the volatility of B2O3 in borosilicate [BS] liquid is low, not all B2O3 could have escaped from the melt. In positive mass spectra, the 27.8 amu atomic mass reflects 11B-16O with intensities that reflect the boron isotope natural abundance. A detailed SIMS analysis was carried out on the oxidised layer. Fig. 3 shows an EI of a region in the oxidised layer 1 obtained at a tilt angle of 45  in the FIB workstation revealing both bright and dark phases. SIMS analysis was done on some grains, which left shaped craters as shown in Fig. 3. SIMS of crater 1 shows the presence of La, Zr, and Zr O in bright grains suggesting them to be lanthanum zirconate formed according to Eq. (1).  2ZrO2(s) + La2O3(s) → La2 Zr2O7(s)  (1)  Longer exposure to sputtering Ga ions deepened the craters and the grains below the bright phase were also exposed (labelled as (2)). The positive mass spectra collected for longer exposure time on region 2 shows the presence of B and Si along with Zr and Zr O. The brightness in the FIB-EI images tends to indicate the electron conductivity. The electron conductivity in zirconate decreases with an increase in lanthanum content [33], which suggests that the dark phase must contain higher amounts of lanthanum (this could not be quantified using present techniques). Also, dark contrast suggests its insulating nature corresponding to a glass composition of La Zr Si O. This suggests the presence of mixed oxide phases comprising silicate glass and crystalline zirconate. To characterise more precisely the crystalline nature of different phases in oxidised layers, they were further analysed using TEM on FIBed sections. Fig. 4(a) shows an electron image of the ion polished region of the  oxidised layer in ZSLO obtained at a tilt angle of 0  in the FIB workstation and Fig. 4 (b) shows the trench of 15 \\u242em × 5 \\u242em made on the rectangular box in Fig. 4 (a) for lifting out the TEM sample at 45  tilt angle. Platinum was deposited on top of the trench to improve the conductivity of the section. The cross-sections observed at 0  and 45  show that the oxidised layer consists of several phases. Fig. 4 (c) shows a bright-field TEM image from the interface region between intermediate layers 2 and 3 in Fig. 2 (a) revealing that the intermediate layers consist of various crystalline and amorphous phases. Boxed regions labelled 1 and 2 in Fig. 4 (c) were further examined. Fig. 5 (a) shows the bright-field TEM image of box 1 in Fig. 4 (c) at higher magnification. Most grains have rounded surfaces indicating that they have been involved in some reactions during oxidation. Three distinct contrast phases were seen: light, medium and dark grey. In some regions, the grains do not have any definite shape and are presumably from liquid formed at high temperature. Selected area electron diffraction [SAED] patterns and EDS were taken from the labelled grains in Fig. 5(a). EDS of grains A, B and C show that they were mainly Si and C with a trace of oxygen. SAED patterns taken on grains A and C at the [11-20] zone axis reveal that they were 6H-SiC. The characteristic streaks observed in the SAED pattern reveal the presence of stacking faults in the grains. The light grey phase D did not have any defined shape and it appeared that it is spread all over and has potentially filled sharp voids. EDS analysis on this phase shows that it is Si and O with a trace of Zr. The black dots appearing in most grains are sputtered Ga. The continuous phase might be silicate phase with some ZrO2 inclusions or  \\x0c', 'D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  4063  (a) shows the electron image of an ion polished region of the oxidised layer in ZSLO obtained at a tilt angle of 0  Fig. 4. in the FIB workstation, (b) shows the trench of 15 \\u242em × 5 \\u242em made on the rectangular box in (a) for lifting out the TEM sample at 45  tilt angle and (c) shows the bright field TEM image of an FIB milled region. Boxes 1 and 2 in (c) are detailed in Figs. 5 and 6.  precipitation matrix.  of  some  ZrSiO4  grains  (Eq.  (2))  in  a  silicate  2SiO2(s) + ZrO2(s) → ZrSiO4(s)  (2)  Precipitation of ZrSiO4 as well as the presence of ZrO2 particles has been detected [32] in the BS layer and on the outer surface of ZrB2 -based UHTCs. Such phases were more evident with longer oxidation times. No crystalline silica was detected in the present study and the glassy character of the BS layer is clearly visible in microstructures of oxide layer. In addition, crystallisation phenomena usually occur under equilibrium conditions, which are not present in this study due to continuous phase separation and precipitation, vapourisation or diffusion of different species during oxidation. Medium grey grains such as H in Fig. 5(a) mainly consist of Zr and O. SAED patterns from these grains can be indexed with monoclinic crystal structure. EDS analysis taken over dark grey E and G shows La, O, Si and a trace of Zr. This confirms the formation of lanthanum monosilicate via reaction (3) and SAED (Fig. 5) confirms its crystalline nature.  SiO2(s) + La2O3(s) → La2 SiO5(s)  (3)  Region F is mainly Si, C and O with minor Zr and a weak La peak. This continuous phase may be mixed phases of SiOxCy , ZrOxCy and  Lax ZryOz . Similar phases in the intermediate layer were observed previously in the ZrB2 -based UHTCs added with La2O3 after oxidation from 0 to 4 h at 1600  C.[13] SiOxCy and ZrOxCy particles were detected within the BS melt after oxidation for 4 h at 1600  C, presumably as a consequence of its higher viscosity from the high La2O3 content (10 weight %). The ZrOxCy and SiOxCy particles appear to contribute some protection by becoming oxidised after longer exposures times, explaining the improved oxidation protective behaviour observed in oxidation kinetics for ZSLO [32]. SiO2 and ZrO2 particles could react with C after SiC oxidation according to Eqs. (4) and (5).  ZrO2(s) + xC → ZrOx Cy  SiO2(s) + xC → SiOx Cy  (4)  (5)  The oxycarbides formed via Eqs. (4) and (5) may have been formed and become fully oxidised at the top (outer) surface after long exposure times [32]. Another possible route for the formation of oxycarbides is by the transport of O2 species through the BS layer, O2(g) from the atmosphere and CO2(g) released from SiC(s) . However, while these oxycarbides remain in the intermediate layers, even after long exposure times they do not oxidise further to form SiO2 and ZrO2 . This observation suggested that oxycarbide  \\x0c', '4064  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  Fig. 5.  (a) A bright-field TEM image of box 1 in (c) at higher magnification. EDS and SAED corresponding to grains labelled are shown.  coatings were suitable candidates for UHTC oxidation protection systems [32]. Fig. 6 is a bright-field TEM image of another region in the oxidised layer (box 2 in Fig. 5(c)). Different features and morphologies are observed in the area analysed compared to box 1. As observed in other regions of the oxidised layer SiC grains were rounded and indexed as cubic 3C polytype from the SAED pattern shown in Fig. 6. Previously we reported the presence of 3C-SiC in hot pressed SiC-ZrB2 samples which used 6H SiC starting powder and suggested a dislocation mechanism may be responsible for the temperature and deformation induced transition [34]. While the polytypic transformation from 6H to 3C SiC could have occurred during pressure assisted sintering due to dislocation mechanisms, it can also occur by the same mechanism as a consequence of  oxidation. Similar contrast phase B is adequately wetting the SiC grains along their surfaces and also starting to fill voids owing to its fluid nature. EDS shows that B is predominantly Si and O suggesting that SiC oxidises first to form SiO2 and subsequently becomes molten at the melting temperature of silica (1600  C) although this may be reduced by the fluxing action of B. The free carbon formed during the initial stage of oxidation should have volatilised as CO2 or reacted to form oxycarbides during the 1 h holding time. EDS taken on region C shows C and O with less intense Si and Zr peaks. Their rounded morphologies in the bright-field TEM image suggest strongly that both B and C were solidified from a melt. Other possible oxidation reactions indicate that other phases such as La2 SiO5 , ZrO2 , SiOxCy , ZrOxCy or ZrSiO4 , exist together in the intermediate layers and which at least partially provide the barrier  \\x0c', 'D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  4065  Fig. 6.  (a) The bright field TEM image from box 2 in Fig. 4(c) at higher magnification. EDS and SAED corresponding to grains labelled are shown.  for O2 diffusion into the bulk. A possible reaction sequence during oxidation is, ZrB2 and SiC oxidise to ZrO2 , B2O3 and SiO2 phases and CO2 gas is emitted. This leads to formation of borosilicate glass. Later ZrO2 and RE2O3 grains might dissolve in this borosilicate glass. At high temperature, B2O3 volatilises leaving behind ZrO2 and RE2O3 grains in a silicate melt and a competition exists between ZrO2 and RE2O3 to react with SiO2 . This leads to the formation of both RE2 Zr2O7 and/or RE2 Si2O7 phase(s) along with unreacted ZrO2 and liquid silicate. These intermediate phases further protect the bulk material lying below. The oxides may either dissolve or remain as solid particles in the BS liquid. Solids may rise due to density difference in the convection currents in the BS melt and dissolved oxides may precipitate on the surface. This may explain why we observe a refractory coating on the surface with an intermediate layer of BS glass (Figs. 2(a) and 7(a)).  3.2.2. Oxidation of ZS20/10 wt.% LaB6 (ZSLB)  (250 \\u242em) ZSLB had similar oxidation layer thickness as observed in ZSLO after oxidation for 1 h at 1600  C [13]. Microstructure and SIMS analyses from different regions in the cross-section of the oxidised surface are shown in Fig. 7. The oxidised surface is heterogeneous and contains a variety of phase morphologies. In some regions, it appears to be compact with grain boundary phases presumably derived from liquid. Two layers can be seen at low magnification in Fig. 7(a). The outer layer is 120 \\u242em thick and dense  while the intermediate layer is 50 \\u242em thick and porous (Fig. 7 (b)). The porosity present in the intermediate layer may reduce strength degradation after oxidation [7]. Our previous work [13] verified that the continuous phase in the outer layer contained mainly Zr, La, and O while the isolated dark regions were mostly silica suggesting that they derive from liquid since silica would melt at this temperature. Positive ion SIMS data was collected from different regions in the cross-section front of the oxidised surface and is shown in Fig. 7(c). Region 1 is La, Zr and their respective oxides (Zr O and La O) indicating that the top surface layer must be either La2 Zr2O7 or mixture of La2O3 and ZrO2 or both. The presence of Si in the SIMS spectra could be due to the overlap of phases present as the layers are ragged. SIMS of region 2 reveals more B and Si along with Zr, La, Zr O and La O. Boron was not observed on the top exposed surface owing to the rapid evaporation of boria from the BS above 1500  C [39]. At 1600  C, BS vaporisation is more rapid than BS formation. However, in the intermediate oxide layer, where temperature is below the boria evaporation temperature, the BS melt mixes more homogeneously with the porous oxide layer as a consequence of increasing BS melt viscosity. This indicates that the layer below the top oxidised layer might consist of lanthanum silicate/lanthanum borate/borosilicate or all along with oxides of La and Zr. Regions of layer 3 and 4 have identical composition with peaks corresponding to B, Si, Zr, La, La O and very weak Zr O. This suggests that some LaB6 [35] and ZrB2 [13] phases have been  \\x0c', '4066  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  (a) SE image of the cross-section of the oxidised surface of ZSLB after oxidation for 1 h at 1600   C showing different oxidised layers, (b) FIB-EI of cross-section showing Fig. 7. different layers and (c) SIMS collected from marked (1-4) regions.  oxidised to respective oxides at temperatures lower than 800  C but as there is no Si O peak, it can be assumed that the temperature was not high enough for the SiC [17] to oxidise with the O2 partial pressure locally present in this region. Fig. 8 shows FIB electron images of a region in the oxidised layer (layer 2 in Fig. 7) revealing phase separation in the liquid. Different contrast phases with distinctive features are seen in Fig. 8(a). The texture of each phase with curved and wavy boundaries reveals they were liquid at high temperature. In Fig. 8(b), the phase in the centre of the image contains isolated dark contrast spheroids and a lighter continuous phase. SIMS analysis was carried out inside the dark spheroids and outside them in the continuous light phase and the corresponding craters are seen in Fig. 8(b). Different levels of B, Si and La are observed in both regions. In particular, La O is more obviously present inside the spheroids. The continuous phase is likely lanthanum borosilicate glass containing some Zr from which an immiscible differing composition La2O3 -containing BS glass also containing Zr has separated presumably with very different viscosity. Fig. 9 shows a FIB-EI from this sample with a region at its centre cleaned with Ga ion at 1.05 nA current. A section of the image has been coloured to clearly reveal the different phases. The purple phase is mainly of the La Zr O system, the red phase consists of La B Si and the dark spheroids inside the red phase contain La O rich B Si; the green regions are pores.  Fig. 10(a) is a bright-field TEM image of the oxidised layer of ZSLB revealing different features labelled A-E which were further analysed by EDS and SAED. Region A is shown in Fig. 10(b) at higher magnification. EDS from A in Fig. 10(c) shows that it contain Zr, La and O and the corresponding SAED pattern (Fig. 10(d)) was from a cubic [0 0 1] zone axis. This confirms that A is La2 Zr2O7 having cubic pyrochlore structure. The bright-field TEM image of grain B at higher magnification is shown in Fig. 10(e) and EDS of grain B in Fig. 10(f) confirms it to be ZrO2 . Another dark phase C consists of La-rich Si, Zr and O and the SAED taken on this grain reveals its non-crystalline nature. However, the EDS of the light grey phase D (Fig. 10(g)) consists of only La, Si, O and trace of Zr. The inset in Fig. 10(g) shows its amorphous nature. Fig. 10(h) is the EDS taken on bright phases E within glass matrix D showing the presence of Si, La and O. The SIMS and TEM-EDS analyses reveal that the top oxide surface consists of La2 Zr2O7 with localised pockets of silicate. The intermediate layers mainly consist of silicate glass containing La and Zr. In some regions, phase separation with La rich silicate glass has been observed. ZrO2 crystals were precipitated within the silicate glass matrix. SIMS analysis clearly shows the presence of boron in the silicate glass, which is consistent with the argument that boria evaporation is slow in the intermediate layers and there is clear evidence of phase separation or immiscible liquids in the intermediate oxidised layer.  \\x0c', 'D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  4067  Fig. 8.  FIB electron images of a region showing immiscible liquid. SIMS collected from two different regions inside and outside the dark sphere are shown.  Fig. 9. FIB electron image of a region in the oxidised layer of ZSLB showing different phases and immiscible liquid. A section of this image has been coloured to clearly show different phases. The purple phase is mainly of the La Zr O system, the red phase consists of La B Si and the dark spheroids inside the red phase contain La O rich B Si; the green regions are pores.  \\x0c', '4068  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  Fig. 10. (a) Bright-field TEM image of a region in the oxidised layer of ZSLB showing different phases, (b) Region A is shown in Fig. 10(a) at higher magnification, (c) EDS from A showing Zr, La and O, (d) SAED pattern of A taken from a cubic [0 0 1] zone axis, (e) The bright-field TEM image of grain B, (f) EDS of grain B showing Zr and O peaks (g).  3.2.3. Oxidation of ZS20/10 wt.% Gd2O3 (ZSGO)  ZSGO after oxidation for 1 h at 1600  C had its top surface covered mostly by crystalline oxides [13] in contrast to the amorphous silicate coating formed in base line SiC-ZrB2 UHTCs.[17,18,36,37]  Fig. 11 (a) shows a BS electron image of cross-section showing different layers with irregular top layer thickness. Fig. 11(b) shows a bright-field TEM image of the oxidised layer of ZSGO. Two different contrast phases predominate: a dark grey phase on the right side  \\x0c', 'D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  4069  Fig. 11. (a) Bright-field image of a region in the oxidised layer of ZSGO showing different phases, (b) Regions near the interface of two grains in (a) are observed at higher magnification. Corresponding SAED patterns and EDS graphs are shown.  and light grey phase on the left side. Regions near the interface of two grains are observed at higher magnification and analysed in Fig. 11(c). The SAED pattern from dark phase ‘Z’ is from monoclinic [0 1 1] confirming its crystalline nature and the EDS shows Zr and O. It can be concluded from this analysis that m-ZrO2 is present close to the top oxide surface. Previous studies on the oxidation of rare earth doped ZrB2 -based UHTCs explained the presence of Re2 Zr2O7 and ZrO2 on the top exposed surface [32]. ZrOxCy particles were also detected by EDS within the BS melt and on top of the oxide layer. This top ZrOxCy eventually also gets fully oxidised with time and forms ZrO2 , while intermediate ZrOxCy still remains deeper in the oxidised layer after long exposures times. Similar behaviour  was found for formation of SiOxCy particles. Therefore, the ZrOxCy and SiOxCy particles formed contribute to oxidation protection at later times in the oxidation process, which could account for the improved oxidation protective behaviour observed in ZSLO over extended times. The ZrO2 particles could react with C(s) after SiC oxidation producing ZrOxCy . The mechanism according to this reaction is that ZrOxCy particles are formed on the oxide layer interface and become oxidised on the top surface after long exposure times. The light grey phase in Fig. 11(b) contains homogeneous distribution of dark spheroids. EDS on these submicron spheroidal particles reveals predominantly Gd, Si and O with a trace of Zr. The light grey continuous phase contains more Si and O with smaller amounts of  \\x0c', '4070  D.D. Jayaseelan et al. / Journal of the European Ceramic Society 35 (2015) 4059-4071  Gd and Zr. EDS in the TEM utilised an ultrathin window and the semi quantitative analysis package so that accurate chemical compositions could not be determined. SAED pattern (Fig. 11) confirms both phases are glassy but phase separated. EDS indicates that the dark isolated glass is Gd rich compared to the light continuous glass. Such phase separation has also been observed in the oxidised layer of ZSLB in the present study and previously by other researchers [38,39]. The observation of phase separation is important since it is known to increase glass viscosity leading to slower oxidation in such systems, In all three samples (ZSLO (Fig. 2(a)), ZSLB (Fig. 7(a)) and ZSGO after oxidation for 1 h at 1600  C d, (Fig. 11(a))) the common features observed after oxidation at 1600  C from their crosssection microstructures are a protective top crystalline oxide layer, intermediate layer(s) with BS containing immiscible liquid and unoxidised bulk material. These observations have been used to design a candidate oxidation and thermal protection system using UHTCs in which reduced strength degradation after oxidation is expected [40]. Although, ZrB2 -SiC oxidised at 1600  C had much thinner and more compact coatings [13], due to the fact it is a glassy phase, at higher velocity for a long time it becomes soft and will be blown off subsequently. Furthermore, in the presence of water vapour, silica is not an ideal candidate material as it readily reacts with water to form Si(OH)4 which will result in the loss of material [41]. When comparing the cross-section micrographs of ZSLO [13], ZSLB and ZSGO after oxidation at 1600  C, they have similar thickness of oxidation layer. However, the notable difference between them when comparing the chemical composition of the top and intermediate layers is that the presence of boron, which is higher in ZSLB than other two compositions. It can be derived that ZSLO and ZSGO will have edge over ZSLB in terms of protection around 1600  C.  4. Conclusions  Detailed structural and chemical analyses of the oxide layers in rare earth containing ZrB2 -based UHTCs after oxidation for 1 h at 1600  C revealed that rare earth additions aid formation of crystalline high melting refractory oxide(s) such as lanthanum zirconate which, along with ZrO2 and silicate phases were observed on dense, thick outer oxidised surface layers. The layer-by-layer analysis across the cross-section of oxidised samples revealed that boron remained even on the top (outer) surface. However, the amount of boron present in the deeper intermediate layers is higher most likely due to boria evaporation from the surface. B loss would increase the viscosity of the silicate in the oxidised layer thereby decreasing the oxygen diffusion. Boron was found predominantly in the oxidised layers of ZSLB, which had more boron due to the addition of LaB6 . The intermediate layers in all three samples contain mixed phases such as silicates and oxycarbides, the latter of which eventually become oxidised likely contributing to a slowing of the oxidation process. FIB electron and bright-field TEM images also showed the existence of phase separation in the glasses formed also increasing glass viscosity aiding oxidation protection.  Acknowledgements  One of the authors, D.D.J. acknowledges the support of DSTL, UK for providing the financial support for this work under contract number DSTLX-1000015413. E.Z.S. acknowledges the support of Fundación Ramón Areces, Spain and the Centre for Advanced Structural Ceramics (CASC) for his postdoctoral fellowship to stay at Imperial College London, UK. E.Z.S. also acknowledges current support through a contract from the JAE-DOC post-doctoral programme of CSIC, Spain, co-funded by the European Social Fund (FSE).  References  and HfB2 -based ultraJ. Mater. Sci. 39 (2004)  [1] E. Opila, S. Levine, J. 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},{
  "_id": 250,
  "PDF": "Structural evolution and ablation mechanism of a hafnium carbide coating on a C-C composite in an oxyacetylene torch environment.pdf",
  "Text": "['Corrosion Science 61 (2012) 156-161  Contents lists available at SciVerse ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Structural evolution and ablation mechanism of a hafnium carbide coating on a C/C composite in an oxyacetylene torch environment  ⇑  Ya-lei Wang, Xiang Xiong  , Xue-jia Zhao, Guo-dong Li, Zhao-ke Chen, Wei Sun  State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, PR China  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 29 December 2011 Accepted 21 April 2012 Available online 30 April 2012  Keywords:  A. Ceramic B. Weight loss C. High temperature corrosion  1. Introduction  HfC coating was deposited on carbon/carbon composites by chemical vapor deposition. The evolution of morphology and microstructure of the coating during ablation is investigated and the ablation mechanism is proposed. Results show that the coating shows an outstanding ablation resistance. After ablation, three distinct regions (central, transitional and outer ablation region) can be detected on the coating surface. In the former two regions, four layers can be divided from the cross-section of the ablated coating. Both molten HfO2 and HfCxOy could act as oxygen diffusion barrier during ablation. Oxidation and mechanical erosion are main ablation behavior of the HfC coating. Ó 2012 Elsevier Ltd. All rights reserved.  In the high-technology field of rocket solid propulsion, thermal protection materials constituting the nozzle and divergent inner walls are submitted to very high temperatures (2800 °C). Under such heavy solicitations, few materials can remain in the solid state [1]. Carbon/carbon (C/C) composites are a family of materials that process extraordinary and unique characteristics that make them attractive for use in aerospace application [2-4]. However, the application of C/C composites is restricted by the poor oxidation resistance of carbon, which readily oxidizes above 500 °C [5-7]. Thus, it is necessary to improve the oxidation resistance of C/C composites in ultrahigh temperatures and oxidizing atmosphere. Deposition of external protective coatings is considered to be one of the most reasonable choices to mitigate effects of oxidizing atmosphere that otherwise limit the performance of C/C composites [6,8-12]. It is well known that refractory metal carbide (e.g., HfC, ZrC, TaC, NbC etc.) can withstand the extreme thermal and chemical environment due to their high melting point, high hardness and retained strength at high temperatures [13,14]. Among these refractory metal carbides, HfC possesses a melting point as high as 3890 °C [15], and is one of the most promising coating materials for C/C composites in ultrahigh temperature applications due to its good high-temperature performance. Moreover, the associated oxide of HfC has a sufficiently high melting point (HfO2: 2810 °C) and a relatively low vapor pressure [16]. These characteristics  ⇑ Corresponding author. Tel./fax: +86 731 88836079. E-mail address: xiong228@sina.com (X. Xiong).  0010-938X/$ see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2012.04.033  for C/C  result in HfC as an attractive protective coating material composites at ultra high temperature applications. Chemical vapor deposition (CVD) is an effective method to prepare HfC coating on C/C composites [17,18]. It can control the compound stoichiometry, chemical purity and microstructure of the coating. More important, it can deposit HfC coating on C/C composites at relatively low temperature. HfC coatings for C/C composites have been prepared by Russia scientists, and it is reported that the coatings exhibit outstanding anti-ablation properties in the ignition experiments of solid rocket motor. Wang and Li [19] developed HfC-based multilayer coating on C/C composites with improved anti-ablation properties. However, there are limited studies reported thus far on the structural evolution and ablation mechanism of HfC coating under oxyacetylene flame. Our previous work [20] has prepared multilayered HfC coating on C/C composites by LPCVD. However, before the practical application of HfC protective coating at ultrahigh temperature, the investigation on the evolution of morphology and microstructure during ablation is far required. In this paper, HfC protective coating was deposited on C/C composites by LPCVD with HfCl4-CH4-H2-Ar system. The ablation behavior of HfC coating is investigated by an oxyacetylene torch. The evolution of surface morphology and microstructure of the HfC coating during ablation is investigated. Moreover, the ablation mechanism is proposed.  2. Experimental  2.1. Coating preparation  HfC coating was deposited in a vertical hot-wall CVD reactor (u30 mm \\x02 10 mm) under reduced pressure (5 kPa). Cylindrical  \\x0c', 'Y.-l. Wang et al. / Corrosion Science 61 (2012) 156-161  157  and those of C2H2 were 0.095 MPa and 0.31 L/s, respectively. The distance between the nozzle tip and the surface of the coated sample was 10 mm and the inner diameter of the nozzle tip was 2.0 mm. The coated sample was fixed in a sample holder coupled with cooling system. During the test, the ablation gun was ignited firstly. After the flame was steady, the flame temperature was measured using an optical pyrometer and reached 3000 °C. Finally, the ablation gun was moved vertically to the surface of the samples for 60 s. The mass ablation rates were determined by weight changes of the samples before and after ablation. The dimension changes of the samples before and after ablation test were measured to obtain the linear ablation rates.  2.3. Coating characterization  The surface morphologies and microstructures of the HfC coating before and after ablation were examined by a scanning electron microscope (SEM, FEI CO., NOVA Nano230). The elemental composition of the ablated coating was identified with electron probe microanalysis (EPMA, JEOL CO., Jxa8230). The phase compositions of the coating before and after ablation were characterized by a D/max 2550 VB + 18 kW rotating target X-ray diffraction (XRD) analyzer (Rigaku Ltd., Japan, Cu Ka radiation).  3. Results and discussion  3.1. Microstructures of HfC coating  Fig. 1 shows the XRD patterns of the as-prepared coating. It shows that the as-prepared coating is composed of cubic HfC corresponding to the standard JCPDS cards (65-6663). Fig. 2 shows the microstructures of the HfC coating. It can be seen that the coating is smooth without obvious cracks (Fig. 2a). The high magnification image (Fig. 2b) shows that the coating surface is closely stacked by  Fig. 1. XRD patterns of the as-prepared coating.  C/C composites with a density of 1.70 g/cm3 were used as the substrates. As for the CVD HfC process, hafnium chloride (HfCl4) powder, CH4, H2 and Ar were used as source materials. A special powder feeder was used to provide a constant supply of HfCl4 powder to the reactor [20]. The delivery rate of the HfCl4 powder was 1.15 g/min, and the gas flow rates of CH4, H2 and Ar were 160, 800 and 800 ml/min, respectively. The deposition temperature and time were 1500 °C and 3 h, respectively.  2.2. Ablation experiment  The ablation test was carried out in an oxyacetylene torch environment. The pressure and flux of O2 were 0.4 MPa and 0.42 L/s,  Fig. 2. Microstructures of the HfC coating deposited on C/C composites: (a) (b) surface SEM image; (c) (d) cross-sectional SEM image.  \\x0c', '158  Y.-l. Wang et al. / Corrosion Science 61 (2012) 156-161  spherical grains, with grain size in a range of 100-200 nm. The HfC coating has a thickness of about 104 lm and has a good adhesion to C/C composites (Fig. 2c). No obvious interlaminar and impenetrate cracks can be observed. From the high magnification crosssectional SEM image (Fig. 2d), the dense HfC coating consists of granular particles about nano-sized.  3.2. Structural evolution  Fig. 3 shows the photograph of the coated C/C composites after ablation for 60 s. It can be found through the macroscopic observation that the coating keeps the integrity and good adhesion to C/C composites, even if the fast heating and cooling during ablation. After ablation, three distinct ablation regions (e.g., central ablation region, transitional ablation region and outer ablation region) can be distinguished on the surface of the ablated sample. The central ablation region covers a white coating without ablation pits, showing an outstanding ablation resistance. The transitional ablation region covers a mazarine ring with staggered lines, while the coating exhibits a light-blue surface in the outer ablation region. Fig. 4 shows the XRD patterns of the HfC coating after ablation for 60 s. It is apparent that the ablation product is monoclinic HfO2 corresponding to the standard JCPDS cards (74-1506), showing the oxidation of HfC coating. The HfO2 has low evaporation and high melting point [21]. It can endure high-temperature ablation in oxyacetylene torch environment. Fig. 5 shows the microstructures of the ablated coating in different ablation regions. These regions exhibit diverse microstructures. In the central ablation region (Fig. 5a), a continuous HfO2 layer is formed on the coating surface with lots of micro-pores distributed uniformly. These micro-pores can be closely related to the release of gaseous oxidation by-product (CO). The HfO2 layer is composed of HfO2 grains, with grain size in a range of 4-8 lm. Moreover, no obvious boundaries appear among these HfO2 grains. The microcracks generated on the coating surface may attributed to the quick cooling from 3000 °C to room temperature (Fig. 5a). The formation of the HfO2 layer is very effective in limiting the inward diffusion of oxygen to the inner coating, enhancing the oxidation resistance of the HfC coating. In the transitional ablation region (Fig. 5b), the coating surface is covered by HfO2 grains. The grain size is in a range of 200-1000 nm, which is much smaller than those in the central ablation region. Moreover, there are obvious boundaries among these HfO2 grains and only sporadic distribution of micro-pores on the ablated coating surface, indicating a  Fig. 3. Photograph of the HfC coated C/C composites after ablation for 60 s.  Fig. 4. XRD patterns of the HfC coating after ablation for 60 s.  mild ablation environment in this region. When it comes to the outer ablation region (Fig. 5c), the coating surface is closely stacked by fine HfO2 grains, with the grain size in a range of 100-250 nm, which is similar to the HfC grains before ablation (Fig. 2b). That is due to the fact that there is no melting phenomenon of HfC grains happened in this region. The discrepancy in grain size of HfO2 in different ablation regions is contributed to the different sinter degree [20]. The cross-sectional backscattered electron (BSE) images of the ablated coating in different ablation regions are shown in Fig. 6. It can be seen that the coating structures exhibit a clear change. The cross-sections of the ablated coating in all ablation regions are composed of a three-layer structure from the coating surface to the substrate/coating interface: outer oxide layer, transitional interlayer and residual carbide layer. The outer oxide layer that formed on the top of the ablated coating is permeated with small pores, and the distribution of these pores varied across the thickness of the outer oxide layer (Fig. 6a). The size and quantity of these pores increase with the distance away from the transitional interlayer/carbide interface. This porous structure, on one hand, provides diffusion channels for the oxidized gases to the inner coating, which debases the oxidation resistance of the coating. However, on the other hand, it could release the high vapor pressure of CO in the carbide/oxide interface, preventing catastrophic damage of the coating. The transitional interlayer, which lies between the outer oxide layer and the residual carbide, exhibits a color of dark-gray and no obvious pores can be detected. The thickness of the interlayer is about 4 lm in the central ablation region (Fig. 6a), and which increases significantly from the central to the outer ablation region (Fig. 6b and c). Meanwhile, the thickness of the porous oxide has a reversed trend. The residual carbide, which exhibits irradiated color, still has a good adhesion to the C/C substrates. Moreover, no obvious boundaries can be observed either between the outer oxide layer and the transitional interlayer or between the transitional interlayer and the residual carbide layer. For better identify the compositions of the transitional interlayer, 16 areas of the cross-section of the ablated coating in the central ablation region were analyzed with electron probe microanalysis. As shown in Fig. 7a, seven areas were analyzed in the transitional interlayer, while four and five areas in the residual carbide layer and outer oxide layer, respectively. Fig. 7 shows the atom ratios obtained from each area as a function of distance from the substrate/coating interface. It can be seen that the C(K)/O(K) ratio decreases rapidly once enter into the transitional interlayer,  \\x0c', 'Y.-l. Wang et al. / Corrosion Science 61 (2012) 156-161  159  (a)   crack  (b)  (c)  pores  pores  µ  10 m  µ  1 m  µ  1 m  Fig. 5. Microstructures of the ablated coating in different ablation regions: (a) central ablation region; (b) transitional ablation region; (c) outer ablation region.  (a)   carbide  outer oxide  (b)  (c)  interlayer  interlayer  interlayer  1   4   12   16   30 mµ  substrate  30 m  µ  30 m  µ  Fig. 6. Cross-sectional BSE images of the ablated coating: (a) central ablation region; (b) transitional ablation region; (c) outer ablation region.  has a remarkably low diffusion coefficient of oxygen at high temperature [23] and can also act as an oxygen diffusion barrier. In addition, electronic defects with high concentration exist in this HfCxOy phase, which results in the different color of the interlayer compared with the outer oxide layer in the BSE images (Fig. 6) [24]. According to the phase transition temperature (1600 °C) of HfO2 from a monoclinic to a tetragonal phase [22], it can be inferred the HfCxOy is a monoclinic phase in the room temperature after ablation at 3000 °C.  3.3. Ablation properties  The ablation properties of the uncoated and coated C/C composites are listed in Table 1. The scatter bands are standard deviations which are calculated based on experimental results of five separate samples. It is clear that with the protective coating, the ablation resistance of C/C composites is significantly improved. After ablation in oxyacetylene flame for 60 s, the ablation rates of the coated C/C composites decreases by an order of magnitude compared to the pure C/C composites. The mass and linear ablation rates of \\x000.05 mg cm\\x002 s\\x001 the coated C/C composites are only and \\x000.86 lm s\\x001, respectively. There are no obvious mass and dimension changes of the coated C/C composites after ablation for 60 s, exhibiting a good configurational stability. Furthermore, it is noticed that both mass and linear ablation rates exhibit negative values, resulting from the weight gain and volume expansion during  Table 1 Mass and linear ablation rates of C/C samples with and without HfC coating. The scatter bands are standard deviations calculated based on experimental results of five separate samples.  Samples  C/C Coated C/C  Mass ablation rate (mg cm\\x002 s\\x001)  1.08 ± 0.12 \\x000.05 ± 0.01  Linear ablation rate (lm s\\x001)  22.79 ± 0.08 \\x000.86 ± 0.02  Fig. 7. Atom ratios obtained from each of the 16 areas as a function of distance from the substrate/coating interface.  and then keeps constant through the width of the transitional interlayer. After the transitional interlayer is crossed, the C(K)/ O(K) ratio falls off again in the outer oxide layer. For the C(K)/ Hf(L) ratio, it shows a similar trend as C(K)/O(K) ratio. In addition, the O(K)/Hf(L) ratios are rather low in the carbide layer. There is 4-7 atomic % oxygen containing in the residual carbide layer, indicating the diffusion of oxygen inwardly. As entering into the transitional interlayer, the O(K)/Hf(L) ratio increases rapidly. This ratio is also constant within the transitional interlayer, and then increases again in the interlayer/oxide interface. It is worth noting that the atom ratios, including C(K)/O(K), C(K)/Hf(L) and O(K)/ Hf(L) ratios, are all constant through the width of the transitional interlayer, suggesting that the composition of this interlayer is constant throughout. Bargeron and Benson [22] have confirmed that it is an oxygen-deficient HfO2 (HfCxOy). The HfCxOy phase  \\x0c', '160  Y.-l. Wang et al. / Corrosion Science 61 (2012) 156-161  Fig. 8. Schematic diagram of the ablated coating in the central ablation region.  the oxidation of HfC, which conceals the mass and linear loss of the coated C/C composites.  3.4. Ablation mechanism  It is well known that the ablation under the oxyacetylene flame is a complex process with thermal chemical, physical and mechanical actions [25-27]. In our present work, the central ablation region suffers the highest temperature and pressure in oxyacetylene flame. There are severe reactions between the reactive gases (mainly O2 and H2O) and the HfC coating. In the present case, the main expected reactions during ablation are as follows:  HfCðsÞ þ O2 ðgÞ ! HfO2 ðsÞ þ COðgÞ  HfCðsÞ þ H2OðgÞ ! HfO2 ðsÞ þ COðgÞ þ H2 ðgÞ  HfO2 ðsÞ ! HfO2 ðlÞ  ð1Þ  ð2Þ  ð3Þ  In the initial stage of ablation, the HfC grain boundaries are oxidized preferentially in the flowing gas to form HfO2 (Eqs. (1), (2)). Thus, a porous oxide layer is formed on the coating surface. In this step, the ablation rates are controlled by the oxidation rate of HfC. During ablation, the central ablation region of the coated sample has the highest temperature of about 3000 °C, which is higher than the melting point of HfO2. So, a molten HfO2 layer forms (Eq. (3)) and spread on the ablation surface in the central ablation region. On the one hand, the molten HfO2 layer partially seals the pores (Fig. 5a) and acts as an efficacious oxidation protective barrier. Thus, the further oxidation of the HfC can only be achieved by the solution and bulk diffusion of oxygen through the dense HfO2 layer. On the other hand, it resists the ultrahigh temperature scouring of oxyacetylene flame due to its high viscosity. It can be seen that the surface of the ablated coating is rather smooth, which is contributed to the scouring of the gas flow with high-temperature and high-velocity. The formation of the molten HfO2 layer with high melting point and viscosity sufficiently improves the ablation resistance of the coating. Oxidation and mechanical erosion are the main ablation behavior of HfC coating in the central ablation region. In the transitional ablation region, the coating surface suffers mild ablation compared with that in the central ablation region due to the lower surface temperature. However, the evolution of the ablated surface in this region is similar as that in the central ablation region (Fig. 5b). Thus, oxidation and mechanical erosion are also the main ablation behavior in the transitional ablation region. While in the outer ablation region, the surface temperature and pressure is far lower than that in other ablation regions, the  oxides on the coating surface are difficult to melt and be blow away. So, oxidation is the main ablation behavior in the outer ablation region. According to the former analyzes, the ablated coating could be divided into four layers both in the central and transitional ablation regions, while three layers in the outer ablation region due to no formation of the molten HfO2 layer. Fig. 8 shows the structures of the ablated coating in the central ablation region. It is clear that the structure of the ablated coating consists of four distinct layers (HfC layer, HfCxOy interlayer, porous HfO2 layer and molten HfO2 layer) from the substrate/coating interface to the top of the ablated coating (Fig. 8). During ablation, both molten HfO2 and HfCxOy could act as an oxygen diffusion barrier to retard the further oxidation of the inner coating. In addition, the molten HfO2 layer is a good thermal barrier, which could hinder the heat transfer and gas flow, reducing the heat attack to the C/C composites. Moreover, large amount of energy from the oxyacetylene flame can be consumed by the oxidation of HfC and the melting/evaporation of the oxide, which reduces the surface temperature and improves the ablation resistance of C/C composites.  4. Conclusions  HfC protective coating was deposited on C/C composites by LPCVD. The ablation properties of the coated C/C samples were tested in an oxyacetylene torch. The evolution of morphology and microstructure during ablation is investigated using various experimental techniques. After ablation for 60 s, three distinct regions (central, transitional and outer ablation region) can be observed on the ablated coating surface. In the former two regions, molten HfO2 layer is formed on the coating surface, which did not happen in the outer ablation region. During ablation, an oxygen-deficient HfO2 (HfCxOy) phase with low diffusion coefficient of oxygen is formed between the residual carbide layer and the outer oxide layer, which could retard the further oxidation of the residual carbide. After ablation for 60 s, the mass and linear ablation rates of the (\\x000.05 mg cm\\x002 s\\x001 and \\x000.86 lm s\\x001, coated C/C composites respectively) are far lower than those of pure C/C composites. Oxidation and mechanical erosion are the main ablation behavior of HfC coating. Both molten HfO2 and HfCxOy could act as an oxygen diffusion barrier to retard the further oxidation of inner coating during ablation. The oxidation of HfC and the melting/evaporation of the oxide could absorb a great amount of heat and reduce the erosive attack on C/C composites.  \\x0c', 'Y.-l. Wang et al. / Corrosion Science 61 (2012) 156-161  161  Acknowledgments  This work is supported by the National Basic Research Program of China (No. 2011CB605805), Creative Research Group of National Natural Science Foundation of China under the Grant No. 51021063, Hunan Provincial Innovation Foundation for Postgraduate and by doctoral Fund of new teachers from Ministry of Education under grant number 20100162120002.  References  [1] G.L. Vignoles, Y. Aspa, M. Quintard, Modelling of carbon-carbon composite ablation in rocket nozzles, Compos. Sci. Technol. 70 (2010) 1303-1311. [2] X.T. Shen, K.Z. Li, H.J. Li, H.Y. Du, W.F. Cao, F.T. Lan, Microstructure and ablation properties of zirconium carbide doped carbon/carbon composites, Carbon 48 (2010) 344-351. [3] F. Smeacetto, M. Salvo, M. Ferraris, Oxidation protective multilayer coatings for carbon-carbon composites, Carbon 40 (2002) 583-587. [4] X. Xiong, Y.L. Wang, Z.K. Chen, G.D. 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Long, Texture structure and ablation behavior of TaC coating on carbon/carbon composites, Appl. Surf. Sci. 257 (2010) 656-661. [10] Q.M. Liu, L.T. Zhang, F.G. Jiang, J. Liu, L.F. Cheng, H. Li, Y.G. Wang, Laser ablation behaviors of SiC-ZrC coated carbon/carbon composites, Surf. Coat. Technol. 205 (2011) 4299-4303. [11] X.T. Shen, K.Z. Li, H.J. Li, Q.G. Fu, S.P. Li, F. Deng, The effect of zirconium carbide on ablation of carbon/carbon composites under an oxyacetylene flame, Corros. Sci. 53 (2011) 105-112.  J. Mater. Sci. 39  the Al2O3-HfO2-  [12] G. Wen, S.H. Sui, L. Song, X.Y. Wang, L. Xia, Formation of ZrC ablation protective coatings on carbon material by tungsten inert gas cladding technique, Corros. Sci. 52 (2010) 3018-3022. [13] S.N. Karlsdottir, J.M. Halloran, Rapid oxidation characterization of ultrahigh temperature ceramics, J. Am. Ceram. Soc. 90 (2007) 3233-3238. [14] Z.Q. Li, H.J. Li, S.Y. Zhang, J. Wang, W. Li, F.J. Sun, Effect of reaction infiltration temperature on the ablation properties of 2D C/C-SiC-ZrC composites, Corros. Sci. 58 (2012) 12-19. [15] A. Sayir, Carbon fiber reinforced hafnium carbide composite, (2004) 5995-6003. [16] S.M. Lakiza, J.S. Tyschenko, L.M. Lopato, Phase diagram of Y2O3 system, J. Eur. Ceram. Soc. 31 (2011) 1285-1291. [17] H.F. Ache, J. Goschnick, M. Sommer, G. Emig, G. Schoch, O. Wormer, Chemical vapor deposition of hafnium carbide and characterization of the deposited layers by secondary-neutral mass spectrometry [J], Thin Solid Films 241 (1993) 356-360. [18] G. Emig, G. Schoch, O. Wormer, Chemical vapor deposition of hafnium carbide and hafnium nitride, J. Phys. IV 3 (1993) 535-540. [19] Y.J. Wang, H.J. Li, Q.G. Fu, H. Wu, D.J. Yao, B.B. Wei, Ablation property of HfCbased multilayer coating for C/C composites under oxy-acetylene torch, Appl. Surf. Sci. 257 (2011) 4760-4763. [20] Y.L. Wang, X. Xiong, G.D. Li, X.J. Zhao, Z.K. Chen, W. Sun, Microstructure and ablation behavior of hafnium carbide coating for carbon/carbon composites, Surf. Coat. Technol. 206 (2012) 2825-2832. [21] S.F. Tang, J.Y. Deng, S.J. Wang, W.C. Liu, K. Yang, Ablation behaviors of ultra-high temperature ceramic composites, Mater. Sci. Eng A 465 (2007) 1-7. [22] C.B. Bargeron, R.C. Benson, X-ray microanalysis of a hafnium carbide film oxidized at high temperature, Surf. Coat. Technol. 36 (1988) 111-115. [23] G.R. Holcomb, G.R. St. Pierre, Application of a counter-current gaseous diffusion model to the oxidation of hafnium carbide at 1200 to 1530 °C, Oxid. Met. 40 (1/2) (1993) 109-118. [24] G.R. Holcomb, The high temperature oxidation of hafnium carbide [D], The Ohio State University, Ohio, 1988. [25] D. Zhao, C.R. Zhang, H.F. Hu, Y.D. Zhang, Ablation behavior and mechanism of 3D/ZrC composite in oxyacetylene torch environment, Compos. Sci. Technol. 71 (2011) 1392-1396. [26] Z.Q. Li, H.J. Li, W. Li, J. Wang, S.Y. Zhang, J. Guo, Preparation and ablation properties of ZrC-SiC coating for carbon/carbon composites by solid phase infiltration, Appl. Surf. Sci. 258 (2011) 565-571. [27] Z.K. Chen, X. Xiong, G.D. Li, Y.L. Wang, Ablation behaviors of carbon/carbon composites with C-SiC-TaC multi-interlayers, Appl. Surf. Sci. 255 (2009) 9217-9223.  \\x0c']"
},{
  "_id": 251,
  "PDF": "Structure evolution of ZrB2–SiC during the oxidation in air.pdf",
  "Text": "['Structure evolution of ZrB2-SiC during the oxidation in air  Xing-Hong Zhang,a) Ping Hu,b) and Jie-Cai Han  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, People ’s Republic of China  (Received 22 November 2007; accepted 2 April 2008)  The structure evolution and oxidation behavior of ZrB2-SiC composites in air from room temperature to ultrahigh temperature were investigated using furnace testing, arc jet testing, and thermal gravimetric analysis (TGA). The oxide structure changed with the increasing temperature. SiC content has no apparent influence on the evolution of structure during the oxidation of ZrB2 -SiC below 1600 °C. However, the evolution of structure for ZrB2 -SiC above 1800 °C was significantly affected by the SiC content. The formation of the SiC depleted layer in the ZrB2 -SiC system not only depends on the surrounding conditions of pressure and temperature but also on the structure distribution of the SiC in the ZrB2 matrix. The apparent recrystallization of the ZrO2 occurred above 1800 °C. The SiC content should be controlled at 16% in the ZrB2 -SiC system for the ultrahigh-temperature application. The mechanisms of structure evolution during oxidation in air were also analyzed.  the  I.  INTRODUCTION  Recent interest in materials for thermal protection systems (TPS) for hypersonic aerospace vehicles and reusable atmospheric re-entry vehicles has resulted in significant research activity focused on ultrahigh-temperature ceramics (UHTC).1-10 Ceramic compounds of refractory borides, carbides, and nitrides such as TaC, ZrB2, ZrC, HfB2, HfC, and HfN belong to the group of materials known as UHTC. In addition to high melting temperatures, ZrB2 and HfB2 have a unique combination of chemical stability, high electrical and thermal conductivities, and resistance to erosion/corrosion that makes them suitable for the extreme chemical and thermal environments associated with hypersonic flight, atmospheric re-entry, and rocket propulsion.5-12 Oxidation resistance is a well-known bottleneck in the development of UHTC for aeropropulsion and hypersonic flight applications. Great efforts have been devoted to this aspect and much progress has been made recently.13-19 Heating ZrB2 in air produces a scale composed of ZrO2 and B2O3. Below about 1100 °C, liquid B2O3 forms a continuous layer that wets the ZrO2 and the underlying ZrB2. In this temperature regime, parabolic (diffusion controlled) kinetics are observed. At intermediate temperatures (1100-1400 °C), the rates of forma Address all correspondence to these authors. a)e-mail: zhangxh@hit.edu.cn b)e-mail: huping@hit.edu.cn DOI: 10.1557/JMR.2008.0251  tion and volatilization of B2O3 are similar, resulting in paralinear kinetics because of competition between mass gain (ZrO2 and B2O3 formation) and mass loss (B2O3 vaporization). Above 1400 °C, active oxidation with rapid linear kinetics has been attributed to loss of B2O3 by evaporation, which leaves behind a porous, nonprotective ZrO2 layer.2,9,11,13,18,19 At 1100 °C and above, the oxidation resistance of ZrB2 can be improved by adding SiC to promote the formation of a silica-rich scale. The silica-rich layer provides protective behavior, which results in mass gain with parabolic kinetics from room temperature up to at least 1600 °C. The addition of SiC not only extends the temperature range of protective behavior, but also imparts the ability to rapidly regain protective behavior after the loss of protection due to excessive temperature, removal of oxide by shear forces, or other causes. Besides SiC, additives such as MoSi2, tantalum compounds, ZrSi2, and other diborides improve the oxidation resistance of diborides either alone or in combination with SiC.13,20-26 In particular, the addition of Ta compounds has been shown to significantly improve the oxidation resistance of ZrB2 -SiC at 1627 °C, whereas it is detrimental to the performance of ZrB2-SiC above 1800 °C. 22,24 In accordance to the previously reported results, a diboride matrix composite including only SiC as a second phase behaves as one of the most promising compositions. The oxidation behavior and mechanism of ZrB2-based UHTC largely depend on the structure evolution. However, to date, the structure evolution of these materials during oxidation in air has not been thoroughly explored, especially for temperatures above 1800 °C.  J. Mater. Res., Vol. 23, No. 7, Jul 2008  © 2008 Materials Research Society  1961  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  The purpose of this work was to investigate the evolution of structure during the oxidation of ZrB2 - SiC composites in air from room temperature to ultrahigh temperature. The mechanism of the structure evolution is also discussed.  II. EXPERIMENTAL PROCEDURE  The samples used here for oxidation testing were fabricated from commercial ZrB2 (Northwest Institute for Non-ferrous Metal Research, Xi’ an, China) and SiC (Weifang Kaihua Micro-powder Co., Ltd., Shandong, China) powders. The ZrB2 and SiC powders had the same purity of 99.5%, with mean particle size of 5 and 2 \\u242em, respectively. The powder mixtures of ZrB2 + 10 vol% SiC, ZrB2 + 20 vol% SiC, and ZrB2 + 30 vol% SiC were ballmilled in ethanol for 8 h and dried in a rotating evaporator. Milled powders were then uniaxially hot pressed in a boron nitride coated graphite die at 2000 °C for 60 min under vacuum and 30 MPa of applied pressure. The heating schedule has been described in detail elsewhere.15 Bulk density and theoretical density were evaluated using the Archimedes method and the rule-of-mixtures, respectively. Coupons were ultrasonically cleaned successively in detergent, deionized water, acetone, and alcohol prior to exposure. The experimental portion of this study focused on exposing ZrB2 -SiC specimens to air at temperatures from 700 to 2000 °C with a step of 100 °C using furnace testing and at the temperatures above 2000 °C using arc jet testing. The oxidation behavior of ZrB2 -20 vol% SiC was also studied using thermal gravi metric analysis (TGA). The weight change was measured under flowing air at a ramp rate of 5 ° C/min up to 1450 °C without an isothermal hold. X-ray diffraction (XRD; Rigaku, Dmax-rb; Tokyo, Japan) was used to identify oxide phases after exposure. The microstructure of each specimen was characterized using scanning electron microscopy (SEM; FEI Sirion, Eindhoven, The Netherlands) along with energy dispersive spectroscopy (EDS; EDAX , Mahwah , NJ) for chemical analysis. The different oxide layers were also investigated after removing the surface layers by polishing parallel to the original surface. The material removal was monitored using optical microscopy so that the desired region was reached.  III. RESULTS AND DISCUSSION  A. Microstructure at room temperature  The relative densities of the prepared ZrB2 -SiC composites are 100%, 98.2%, and 100% for ZrB2-10 vol% SiC, ZrB2-30 vol% SiC, and ZrB 2-30 vol% SiC, respectively. Figure 1 shows scanning electron micrographs of the po l ished surfaces of the ZrB2-based u l trah ightemperature ceramics including 10, 20, and 30 vol% SiC. The darker phase is SiC, and it appears to be uniformly dispersed in the lighter ZrB2 matrix. The microstructures of the composites are regular and little porosity was observed in the polished surfaces. Based on the high relative density and the lack of any open porosity, porosity should not have a significant effect on the oxidation behavior.  FIG. 1. EM images of the polished surfaces of the ZrB2-SiC ultrahigh-temperature ceramics: (a) ZrB 2-10 vol% SiC, (b) ZrB 2-30 vol% SiC, and (c) ZrB2-30 vol% SiC.  1962  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  B. Structure evolution of ZrB2 -SiC during oxidation  ZrB2 exhibited passive oxidation resistance below 1100 °C and the formed liquid B 2O3 acted as a barrier to oxygen diffusion. Above 1300 °C, B 2O3 becomes nonprotective because of rapid evaporation, and the active oxidation of ZrB2 occurs. The introduction of SiC significantly extends the temperature range of the glass as a primary oxygen barrier and increases the maximum working temperature in oxidizing environments. However, the structure evolution and oxidation behavior of the ZrB2 -SiC composites are still not well understood in the whole service temperature range. To better understand the structure evolution of ZrB2-SiC composites during oxidation, we divided temperature into four temperature regions according to their oxidation behavior.  1.  Low temperature range (700-1200  °C)  Thermodynamically, both ZrB2 and SiC should oxidize when exposed to air. However, the oxidation rates of both species are negligible below about 700 °C. Previous studies have reported that the oxidation of ZrB2 by reaction (1) is much faster than oxidation of SiC between 700 and 1200 °C. Assuming that oxidation of ZrB 2 proceeds stoichiometrically, the reaction should produce (melting temperature 450 °C) molten B2O3 and solid ZrO2. Upon cooling to room temperature, the B2O3 forms an amorphous solid, while the ZrO2 is crystalline.  ZrB2共cr兲 + 5 Ⲑ 2O2共g兲 → ZrO2共cr兲 + B2O3  .  (1)  The amounts of B2O3 and ZrO2 on the surface of the specimen oxidized at 700 °C for 1 h were not sufficient to be observed in polished cross sections. To complement the compositional and structural information, the oxidation behavior of the ZrB2-SiC was examined by TGA up to 1450 °C with the same heating rate (5 °C /min) that was used in the furnace oxidation experiments. The change in the mass as a function of temperature is shown in Fig. 2. The weight started to increase just at 700 °C, which corresponds to the temperature at which ZrB2 is reported to begin oxidizing. The weight gain was consistent with SEM analysis that showed a minimal amount of B2O3 formation for specimens heated in air to 700 °C [Fig. 3(a)]. Between 700 and 1200 °C, the weight increased at a low rate.11,13,18 This is consistent with SEM observations that showed the formation and growth of ZrO2 and a protective layer of B2O3 [Figs. 3(b) and 3(c)]. As the temperature rose above 1200 °C, the specimen weight increased slightly, which was attributed to the rapid volatilization of B2O3. Because the oxidation of SiC is much slower than that of ZrB2 in this temperature regime, the SiC particles did not oxidize appreciably, as evidenced by SiC inclusions present in the scale, as shown in Fig. 4(a). The compo FIG. 2. TGA analysis of ZrB2-20 vol% SiC in air up to 1450  °C.  sition of the oxide scale was examined using EDS mapping [Figs. 4(b) -4(d)], which showed that zirconium and oxygen were present along with silicon, suggesting that the reaction layer was composed of ZrO2 and SiC. The distribution of B was not quantified in the present analysis due to the low sensitivity of EDS to light elements. However, EDS maps showed that the outermost layer contained much O but little Si and Zr, which is consistent with the presence of B2O3. In the temperature range from 700 to 1200 °C, the oxide structure consisted of (i) a B2O3-rich outer layer, (ii) a subscale of ZrO2 that contained unoxidized SiC, and (iii) unaffected ZrB2 -SiC. The main driving force for the outer B2O3 scale to grow is the transport of B2O3 (l) from the oxidized ZrB2 to the surface due to a large volume expansion during oxidation. The other possible driving forces for the transport of B2O3, such as a temperature gradient and a chemical potential gradient, are unlikely since isothermal oxidation is assumed in this study and the concentration of B2O3 in the inner oxide scale is lower than that in the outermost layer. In addition, the vapor pressure is low in this temperature regime. With the assumption that oxidation products have theoretical density, 1 unit volume of ZrB2 upon oxidation produces 1.19 unit volumes of solid oxide (ZrO2) and 2.98 unit volumes of liquid oxide (B2O3). Thus, the volume increase during oxidation compared with the initial unoxidized materials varies from 316.6% for pure ZrB2. The liquid oxide must be squeezed up away from the interface to the surface, resulting in the formation of the B2O3-rich surface layer. Therefore, the mechanism of formation of the outer borate layer is the upwelling of borate liquid, with subsequent lateral flow over the outer surface resulting from volume expansion during oxidation. This outer borate layer is an effective barrier to the transport of oxygen in this temperature regime leading to passive oxidation behavior with parabolic mass gain kinetics.  J. Mater. Res., Vol. 23, No. 7, Jul 2008  1963  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  FIG. 3. SEM micrographs of  the surface of ZrB2-30 vol% SiC oxidized at  (a) 700  °C,  (b) 800  °C, and (c) 1200  °C.  FIG. 4. (a) SEM micrograph and EDS maps for (b) O, (c) Si, and (d) Zr for the reaction layer formed by oxidizing ZrB2-30 vol% SiC at 1200 in air for 1 h.  °C  2. Medium temperature range (1200 -1600  °C)  1200 °C, As ZrB2 -SiC was heated above the composition and structure of the surface layers changed. The dominant chemical processes between 1200 and 1600 °C are expected to be the evaporation of B2O3 [reaction (2)] and oxidation of SiC [reaction (3)].  B2O3  (1) →B 2O3  (g)  ,  SiC共cr兲 + 3 Ⲑ 2O2共g兲 → SiO2共l兲 + CO共g兲  .  (2)  (3)  the temperature approaches 1300 °C, As the vapor pressure of B2O3 increases substantially, leading to its  1964  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  rapid evaporation. Meanwhile, SiC starts to significantly oxidize, resulting in the formation of a continuous surface layer above another oxide layer. The weight increased again at this temperature (Fig. 2), which was primarily attributed to the formation of SiO2. A part of the silica was retained in the scale and filled the pores that were not occupied by ZrO2. The other was transported to the surface, which provides an effective barrier to oxygen diffusion that may result in passive oxidation protection with parabolic mass gain kinetics. The driving force for the transport of the silica is also the volume expansion upon the oxidation of ZrB2-SiC, which is similar to that of the B2O3 at lower temperatures. The layered structure in this temperature range consisted of (i) a SiO2 rich glassy outer layer, (ii) a subscale of ZrO2 that contained some SiO2, and (iii) unaffected ZrB2 -SiC. This layered structure is similar to the structure reported for ZrB2-SiC exposed to air at 1327 °C. 3 For ZrB2 containing 30% SiC, the fraction not occupied by ZrO2 is 16.9% and was filled with 25.6% formed silica. The other 74.4% silica was transported to the surface, resulting in the formation of a surface silica layer with 49.1% thickness of the reacted layer, which is consisted with the experimental results [Fig. 5(b)]. (The volume of B2O3 is not included in the calculations at 1600 °C.) An adherent and continuous glass sealed the porosity and covered the exposed faces encapsulating the oxidized specimen. For this reason, the inward diffusion of oxygen through the porosity or grain boundaries toward internally accessible diboride grains was retarded. Based on isothermal studies, the SiO2-rich glassy layer  remains protective up to at least 1600 °C. Because SiO 2 is significantly less volatile than B2O3 at these temperatures, the SiO2-rich layer provides oxidation protection for ZrB2 -SiC over a much greater temperature range than B2O3 does for pure ZrB2. In the temperature range from 1200 to 1600 °C, the material containing high SiC content is more resistant than that containing low SiC content. ZrB2 and SiC have a similar oxidation rate at this stage. It should be noted that the formed ZrO2 has apparently not changed the initial framework of the ZrB2 below 1600 °C as shown in Figs. 4 and 6. The structure of the subscale is similar to that of the bulk in which only ZrB2 and SiC were replaced by ZrO2 and SiO2, respectively.  3. High temperature range (1600-1800  °C)  SiC content has no apparent influence on the evolution of structure during the oxidation of ZrB2 - SiC below 1600 °C. However, the evolution of structure for ZrB 2 - SiC above 1600 °C was appreciably affected by the SiC content. No SiC depleted layer was observed for the Z rB 2 - S iC spec imen s a f te r ox ida t ion in a i r be low 1600 °C [Figs. 5(a) and 5(b)]. However, the SiC depleted layer was detected in both specimens containing either 20 or 30 vol% SiC after oxidation in air at 1700 °C [Figs. 5(c) and 5(d)], which is consistent with the passive- active transition of SiC in air [reaction (4)].27,28 The previous studies showed that the passive-to-active oxidation of SiC in air occurred in the temperature range 1600-1700 °C. 27,28 Figure 7 shows SEM images at high  FIG. 5. Cross-sectional micrographs of oxidized materials for 1 h: (a) ZrB2-20 vol% SiC, 1600 20 vol% SiC, 1700 °C; and (d) ZrB 2-30 vol% SiC, 1700 °C.  °C; (b) ZrB 2-30 vol% SiC, 1600  °C; (c) ZrB 2-  J. Mater. Res., Vol. 23, No. 7, Jul 2008  1965  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  As the temperature approached 1800 ° C, the oxide structure was significantly changed, and oriented growth occurred, as shown in Figs. 9 and 10. No SiC-depleted layer was detected in ZrB2-10 vol% SiC at this temperature, as shown in Fig. 10. At 1800 °C, this temperature may be favorable for the recrystallization of ZrO2. The oriented growth of the scale was primarily due to the evolution of the gaseous byproducts and flows of the silica-rich liquids, which promoted the growth of the zirconia parallel to the direction of the discharge of the gas products and transport of the liquids. EDS maps for Figs. 11(a) O, 11(b) Si, and 11(c) Zr for the reaction layer formed by oxidizing ZrB2-20 vol% SiC at 1800 °C in air for 1 h indicated that layers containing Zr, O, and Si were formed. The EDS map of silicon in Fig. 11(b) shows a sharp transition at the boundary of the virgin material corresponding to SiC-depleted region. The reduction of Si content in the innermost oxide layer is attributed to active oxidation of SiC. Only minor amounts of the silicon were detected in the SiC-depleted layer. The content of Si in outside scale is richer than the inner scale as shown in Fig. 11(b), which is attributed to the convection of Si liquid flow. As can be seen from Figs. 9(a) and 11(b), the glassy SiO2 silica was distributed at the surface and the inner oxide scale, which plugs pores, covers the sample surface , and efficaciously seals ZrO2 grain boundaries. At this stage, the liquid silica glass acts as an effective obstacle against the inward diffusion of oxygen along short-circuit paths (i.e., residual porosity, cracks and grain boundary), and it is helpful to form a dense coherent oxide scale at the same time. Note that the thickness of the silica-rich layer is apparently thinner than that formed at lower temperatures due to the active oxidation of SiC. The crystalline zirconia exhibited a columnar structure in the oxidized ZrB2 containing 20 vol% SiC, whereas this phenomenon was not observed in ZrB2 containing 30 vol% SiC after oxidation under the same condition. Unfortunately, the growth behavior of zirconium oxide in the ZrB2-SiC system at high temperature has not been reported in the previous literature. Here, we argue that the growth of zirconium oxide was mostly dependent on the fraction and distribution of the zirconium oxide in the oxide scale. With an assumption that ZrB2 was completely transformed into rigid solid ZrO2 and was retained in the initial reacted region, the volume fraction for ZrB2 containing 20 vol% SiC is 95.0%, which makes the most of the ZrO2 particles contacting one another. This will increase the growth rate of the ZrO2 solid grains, resulting in the visibly oriented growth of the ZrO2, as shown in Fig. 9(a). However, the volume fraction of ZrB2 containing 30 vol% SiC is only 83.1%; the other fraction will be filled with the liquid oxides. The increased liquid fraction leads to the increased distance between solid grains of ZrO2, which in turn decreases the  FIG. 6. SEM image at high magnification of the layered structure formed after exposure of ZrB2-30 vol% SiC to air at 1600 °C for 1 h.  magnification and line analysis of the cross-sectional micrograph formed after exposure of ZrB2 -30 vol% SiC to air at 1700 °C. The depleted SiC particle in the crosssectional micrographs is evident.  SiC共cr兲 + O2 共g兲 → SiO共g兲 + CO共g兲  .  (4)  Recently, the formation of SiC-depleted layer has been extensively studied by Fahrenholtz from a thermodynamic point of view.17 In fact, the formation of the SiCdepleted layer in the ZrB2 -SiC system not only depends on the surrounding conditions of pressure and temperature but also on the structure distribution of the SiC in the ZrB2 matrix. The formation of the SiC depleted layer is impossible if the distribution of the SiC particles in ZrB2- SiC composites is discontinuous even if the thermodynamic conditions are favorable to the active oxidation of SiC. The other necessary condition for the formation of the SiC-depleted layer is that SiC particles in the three dimensions should be continuous since the diffusion of oxygen in ZrB2 is very low. Additional support for this argument is offered by experimental results in which no SiC-depleted layer was observed from the oxide scale of the ZrB2 -10 vol% SiC even oxidized at 1900 °C (not shown). Figure 8 shows the morphologies of the SiC-depleted region, which are parallel and perpendicular to the top surface of the oxidized specimens. Note that SiC forms a network interconnected in three dimensions, as shown in Fig. 8, although it is discontinuous in two dimensions, as can be seen in Fig. 1. It should be noted that the thickness of SiC-depleted layer in the high SiC content (30 vol% SiC) is much higher than that in the low SiC content (20 vol% SiC). Apparently, the degree of SiC interconnectivity in the matrix increases with increasing SiC content. A high degree of SiC interconnectivity causes the rapid active oxidation of SiC. Consequently, the thickness of the SiC-depleted layer for high SiC content is much higher than that for low SiC under the same condition, as shown in Figs. 5(c) and 5(d).  1966  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  FIG. 7. (a) SEM image at high magnification and (b) line analysis of the cross-sectional micrograph formed after exposure of ZrB2-30 vol% SiC to air at 1700 °C. A and B indicate the start and end of the line scan, respectively.  FIG. 8. Morphologies of the SiC-depleted region, which are (a) parallel and (b) perpendicular to the top surface of the oxidized ZrB2-30 vol% SiC.  FIG. 9. Cross-sectional micrographs of oxidized materials at 1800 °C for 1 h:  (a) ZrB 2-20 vol% SiC and (b) ZrB 2-30 vol% SiC.  growth rate of the ZrO2 solid grains. The micrograph of the oxide scale for ZrB2 -30 vol% SiC at high magnification is shown in Fig. 12. It can be seen that the growth of the ZrO2 was impeded since the majority of the ZrO2 solid grains were separated by silica-rich glass, which confirms the validity of the above analyses. In fact, some ZrO2 will be transported to the surface and then recrystallized. Figure 13 shows a SEM micrograph of the surface of oxidized ZrB2-10 vol% SiC in air at 1800 °C. The whiskerlike features in white contrast were detected in Fig. 13; these were not reported in the previous literature. EDS analyses show that this oxide mainly contains zirconia (ZrO2).  4. Ultrahigh temperature range (above 1800 °C)  The growth of the ZrO2 in the oxide scale was aggravated when temperature was raised to 1900 °C because the further increased temperature favored recrystallization of the ZrO2 grains. At this temperature, the evolution of the structure was significantly dependent on the SiC content, especially for the evolution of the SiC-depleted layer. A dense adherent oxide scale composed of ZrO2 and SiO2 was formed for ZrB2 -20 vol% SiC, whereas a loose porous structure was generated for ZrB2-30 vol% SiC (Figs. 14 and 15). Moreover, cracks and spallation beneath the surface layer were also observed. The images  J. Mater. Res., Vol. 23, No. 7, Jul 2008  1967  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  FIG. 10. Cross-sectional micrographs of oxidized ZrB2-10 vol% SiC at 1800  °C for 1 h.  FIG. 11. EDS maps for  (a) O,  (b) Si, and (c) Zr for  the reaction layer  formed by oxidizing ZrB2-20 vol% SiC at 1800  °C in air for 1 h.  of the cross sections revealed a strong degradation of ZrB2 -30 vol% SiC relative to ZrB 2-20 vol% SiC. ZrB 2- 20 vol% SiC exhibited superior oxidation resistance at 1900 °C in air for 1 h. However, the oxidation behavior of ZrB2-30 vol% SiC under the same condition demonstrated the unsuitability of this material for ultrahightemperature applications. The failure of the material was attributed to the damage of the SiC-depleted region. Comparison of the oxidation results between these two kinds of materials indicates that SiC content significantly affected the oxide structure and performance of ZrB2- SiC at 1900 °C. Based on relative oxidation rates of ZrB2 and SiC at this temperature, SiC should be consumed first through the rapid active oxidation leaving behind the connected porosity. Most likely, if there is a channel for the transportation of oxide products of ZrB2, then ZrO2 would be unstable at 1900 °C, which is significantly different from  FIG. 12. SEM image at high magnification of the oxide scale formed after oxidation of ZrB2-30 vol% SiC in air at 1800 °C for 1 h.  1968  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  FIG. 13. SEM micrograph of SiC in air at 1800 °C.  the surface of oxidized ZrB2-10 vol%  FIG. 15. SEM images at (a) low and (b) high magnification and EDS analysis of the layer structure formed from ZrB2-30 vol% SiC after oxidizing at 1900 °C in air for 1 h.  FIG. 14. A cross-sectional image of oxidizing ZrB2-20 vol% SiC at 1900  the reaction layer °C in air for 1 h.  formed by  the oxidation behavior of the ZrB2-based UHTC at lower temperatures. In fact, when the system is above the boiling point of B2O3 (1860 °C), the oxidation resistance of ZrB2 decreases remarkably, and the formed oxidation products will not be adherent to the local unaltered ZrB2. The consumption of SiC will leave behind channels for the transportation of the formed ZrO2 if the volume expansion upon conversion of ZrB2 to ZrO2 cannot compensate the volume of the SiC consumption. Moreover, the formation of the high-pressure gaseous products from the oxidation of SiC would also accelerate the transportation of ZrO2, resulting in the generation of discrete structure. The vapor pressures of the predominant gases for ZrB2 -SiC composites were calculated to compile a volatility diagram at 1900 °C (Fig. 16). The calculations were performed using JANAF data,29 and the details of the construction of thermodynamic stability and volatility of diagram for ZrB2-SiC can be found elsewhere.  19,30  FIG. 16. Volatility diagram for ZrB2-SiC at 1900  °C.  In air, B2O3 has a high vapor pressure, and its pressure is 0.24 atm at this temperature. The oxide products of the SiC have the highest vapor pressures (up to 1.3 atm) at this stage, which exceed 1 atm, as shown in Fig. 16. The discharge of the gaseous phase products with high vapor pressure formed from the oxidation of SiC will promote the removal of the ZrO2. SiC exhibited preferential oxidation in the present system of ZrB2 -SiC, resulting in the formation of the layer depleted SiC grains. At the same time, ZrB2 grains in contact with depleted SiC grains were directly exposed to oxygen. Thus, the oxidation occurred from the boundary between ZrB2 and SiC to inner ZrB2 grains. Apparently, if the oxidized regimes of the ZrB2 matrix are interconnected within the SiC-depleted region and their products are transported to the external oxide scale, the unoxidized matrices will not interconnect with each other when the  J. Mater. Res., Vol. 23, No. 7, Jul 2008  1969  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  oxide regimes meet together. Consequently, the scale for this material is visually nonadherent, and some oxide spalled (Fig. 15). On the other hand, the rupture of the scale may not occur when the oxidized regimes of the ZrB2 matrix are not interconnected in the SiC depleted region, as shown in Fig. 14. Whether the oxidized regimes of the ZrB2 matrix in SiC-depleted region are interconnected or not mostly depends on the content of SiC in the matrix. Apparently, the average thickness of ZrB2 matrix between two adjacent SiC particles for high SiC content is smaller than that of low SiC content. Therefore, the oxidized regimes of ZrB2 matrix in high SiC content are liable to meet together in comparison to low SiC content. Obviously, if the SiC particles are not interconnected in two and three dimensions, the SiCdepleted region will not be developed when these materials are exposed to oxidizing environments at high temperature since the diffusion of oxygen in ZrB2 is very low. Nonetheless, if the content of SiC is low enough, the oxide scale will be spalled due to large volume expansion and little liquid for the cementation of solid ZrO2 grains. ZrO2 will be recrystallized into an integrality with the further increase of temperature (i.e., >2000 ° C), which was  not reported in the previous study.15 Above 2300 ° C, the softening and degradation of the oxide scale occurred and the severe oxidation was found. Figure 17 shows the macrographs and micrographs of the ZrB2 - 20 vol% SiC composite after arc jet oxidation at 2300 °C for 600 s. A large amount of the molten oxidation product was blown away during the oxidation process, resulting in a high erosion rate and marked change in the configuration [Figs. 17(a) and 17(b)]. The oxide scale was sintered together, as shown in Figs. 17(c) and 17(d). The softening and degradation of the oxide scale resulting from increased temperatures is a major issue of UHTC for use in extreme environments like those associated with hypersonic flight, atmospheric re-entry, and rocket propulsion. The temperature limit for SiC-reinforced ZrB2based UHTC mostly depends on the softening and degrada t ion of ZrO2-based ox ide sca le , wh ich is a lso controlled by SiC content since the melting point of the ZrO2-based oxide scale is closely connected with the content of the formed SiO2, determined by the SiC content. Above 2500 °C, the formed oxide scale is no longer protective in oxidizing environments, and active oxidation with linear kinetics occurs.  FIG. 17.  (a, b) Macrographs and (c, d) micrographs of  the ZrB2 -20 vol% SiC composite after arc jet oxidation at  2300 °C for 1 h.  1970  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  XRD patterns of the surface of the samples after the 1 h oxidation tests at different temperatures are shown in Fig. 18. XRD of the oxide scale showed monoclinic ZrO2 as the major phase for all temperatures. At 1200 °C, only the monoclinic ZrO2 was found; no extra crystalline secondary phase was revealed. Tetragonal ZrO2 was clearly visible at 1500 °C in the XRD patterns (Fig. 18), and the relative intensity increased with increasing oxidation temperature at the beginning and then decreased with the further increased oxidation temperature. The SiO2 phase was evident between 1700 and 1900 °C. After oxidation at 2300 °C, only monoclinic ZrO 2 was detected. The SiC content and microstructure should be optimized to meet the requirements of the ultrahigh temperature application since the performance of the ZrB2-SiC composites is significantly dependent on the SiC content. In this regard, three factors should be taken into account together at ultrahigh temperature: (i) the degree of the SiC interconnectivity, (ii) the volume change upon the oxidation, and (iii) softening and degradation of the oxide scale. To some extent, the SiC interconnectivity with low degree is good for the oxidation resistance of the ZrB2 -SiC composites, which is mostly dependent on the SiC content. The SiC content with low degree is also beneficial to the increase of the temperature limit resulting from softening and degradation of the oxide scale. The SiC content should be less than 30% according to the experimental results. It is beneficial to the formation of a dense coherent oxide scale at ultrahigh temperature when the volume fraction of the transformed rigid solid was controlled at approximately 100%. We assume that the ZrB2 was completely transformed into solid ZrO2 and was retained in the oxide scale, and other liquids were transported to the surface layer. In fact, some silica will be reserved in the oxide scale (Figs. 9, 11, and 12) and  FIG. 18. XRD patterns of ferent temperatures.  the ZrB2-SiC composites oxidized at dif some ZrO2 will be transported to the surface layer. ZrO2 can be transported directly or through liquid convection [Reactions (5) and (6)].  xZrO2共cr兲 + B2O2共l兲 → xZrO2  ⭈B2O3  (l兲  yZrO2共cr兲 + SiO2共l兲 → yZrO2  ⭈SiO2  (l兲  ,  ,  (5)  (6)  The reserved silica in the oxide scale is beneficial to the coupling of the ZrO2 grains and the formation of the adherent dense oxide scale. The volume fraction of the ZrO2 can be set as 100%, considering the volume of the reserved silica in the inner oxide scale and transported zirconium is very small. Their volumes can partly counteract each other under this condition. Therefore, the SiC content is 15.8%, calculated according to this assumption. Moreover, this SiC content is also expected to satisfy the requirement for the degree of the SiC interconnectivity according to the experimental results. Thus, the SiC content should be controlled at 16% in the ZrB2- SiC system for the ultrahigh temperature application.  IV. CONCLUSIONS  The structure evolution and oxidation behavior of ZrB2 -SiC composites in air from room temperature to 2500 °C were investigated using furnace testing, arc jet testing, and TGA. Below 1200 °C, a protective B 2O3rich scale was observed on the surface. As the temperature approaches 1300 °C, the vapor pressure of B 2O3 increases substantially, leading to its rapid evaporation. Meanwhile, SiC starts to significantly oxidize, resulting in the formation of a continuous surface layer above another oxide layer. The driving force for the liquid flow (i.e., B2O3 and SiO2) is the volume increase upon oxidation of ZrB2-SiC. The SiC content has no apparent influence on the evolution of structure during the oxidation of ZrB2 -SiC below 1600 °C. However, the evolution of structure for ZrB2-SiC above 1800 °C was significantly affected by the SiC content. No SiC-depleted layer was observed for the ZrB2 -SiC specimens after oxidation in air below 1600 °C, whereas a SiC-depleted layer was detected in both specimens containing 20 and 30 vol% SiC after oxidation in air at 1700 °C. The formation of the SiC-depleted layer in the ZrB2-SiC system not only depends on the surrounding conditions of pressure and temperature but also on the structure distribution of the SiC in the ZrB2 matrix. The other necessary condition for the formation of the SiC-depleted layer is that SiC particles in the three dimensions should be continuous. The apparent recrystallization of the ZrO2 occurred above 1800 °C, and its behavior was significantly affected by the SiC content. The SiC content should be controlled at 16% in the ZrB2 -SiC system for the ultrahigh temperature application.  J. Mater. Res., Vol. 23, No. 7, Jul 2008  1971  \\x0c', 'X-H. Zhang et al.: Structure evolution of ZrB2-SiC during the oxidation in air  ACKNOWLEDGMENTS  This work was supported by the National Natural Science Foundation of China (50602010), the Research Fund for the Doctoral Program of Higher Education (20060213031), and the Program for New Century Excellent Talents in University.  REFERENCES  1. K. Upadhya, J.M. Yang, and W.P. Hoffman: Materials for ultrahigh temperature structural applications. Am. Ceram. Soc. Bull. 76, 51 (1997). 2. S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, and J.A. 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Lou: Volatility diagrams for silica, silicon n i t r ide , and s i l icon ca rb ide and the i r app l ica t ion to h igh temperature decomposition and oxidation. J. Am. Ceram. Soc. 73, 2786 (1990).  Jr.: NIST-JANAF Thermochemical Tables, 4th ed.  in High Temperature Corrosion  5Si3. J. Mater. Res. 21,  23.  26.  and Materials  1972  J. Mater. Res., Vol. 23, No. 7, Jul 2008  \\x0c']"
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  "PDF": "Study on High Temperature Oxidation Characteristics of NbTi Matrix Composites Reinforced with ZrB2 Particles.pdf",
  "Text": "['Materials at High Temperatures  ISSN: (Print) (Online) Journal homepage: https://www.tandfonline.com/loi/ymht20  Study on High Temperature Oxidation Characteristics of NbTi Matrix Composites Reinforced with ZrB2 Particles  Zongde Liu , Yuan Gao , Ting Wang , Youmei Sun , Yan Gong & Yao Kong  To cite this article: Zongde Liu , Yuan Gao , Ting Wang , Youmei Sun , Yan Gong & Yao Kong (2021) Study on High Temperature Oxidation Characteristics of NbTi Matrix Composites Reinforced with ZrB2 Particles, Materials at High Temperatures, 38:1, 73-81, DOI: 10.1080/09603409.2020.1859814  To link to this article:  https://doi.org/10.1080/09603409.2020.1859814  Published online: 11 Dec 2020.  Submit your article to this journal   Article views: 8  View related articles   View Crossmark data  Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=ymht20  \\x0c', 'MATERIALS AT HIGH TEMPERATURES                 2021, VOL. 38, NO. 1, 73-81  https://doi.org/10.1080/09603409.2020.1859814  Study on High Temperature Oxidation Characteristics of NbTi Matrix  Composites Reinforced with ZrB2 Particles  Zongde Liu  , Yuan Gao  , Ting Wang, Youmei Sun, Yan Gong and Yao Kong  Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of Education, North China Electric Power  University, Beijing, China  ABSTRACT  The oxidation behavior of NbTi matrix composite reinforced by ZrB2 particles at 800°C-1200°C  were investigated. The oxidation kinetic curves of five specimens with different ZrB2 contents  (0-60 wt. %) were obtained, and the phase and morphology of oxidized specimens were  analyzed by XRD, SEM and EDS. The result showed that both temperature and ZrB2 content  have great influence on oxidation resistance of the NbTi-matrix composites. The main oxide  was  TiNb2O7  when  ZrB2  content  was  15%  and  30%.  As  the  ZrB2  content  increased, Nb2Zr6O17 was the predominant oxide. At 800°C and 1000°C, oxide rate reduced  due to the formation of Nb2Zr6O17. When temperature reached 1200°C, the volatilization  of B2O3 caused many cracks and holes and weakened the protective effect of Nb2Zr6O17.  ARTICLE HISTORY   Received 2 July 2018   Accepted 29 November 2020   KEYWORDS   NbTi; ZrB2; high-temperature  oxidation; microstructure  Introduction  Niobium alloy is a kind of good high-temperature  refractory alloy, the use temperature of it is higher  than the traditional Ni-based and Co-based alloy. Its  mechanical properties are good at high temperatures,  the strength can be maintained to 1650°C. It can be  applied to 1100°C at present, and its application at  1200°C or 1500°C is also investigated [1-3]. Nb alloy is  an important high-temperature structural material in  nuclear industry and aviation. However, its oxidation  resistance is poor, and pure Niobium has ‘pest’ oxidation at 600°C, which restricts the application of Nb  and Nb alloys [4]. Many ultra-high temperature materials can improve the oxidation resistance effectively  by adding alloying elements and reinforced particles.  Geng’s studies had shown that the oxidation resistance  of Nb-based alloy could be improved apparently by  adding Ti alloy elements at 800°C and 1200°C [5].  Murayama  investigated  that  if Ti  elements were  added to Nb-Si-Al alloys, the microstructure of Nb Si-Al-Ti alloy was composed of three phases (Nbss,  Nb3Al, Nb5Si3), and the high-temperature strength,  fracture  toughness and oxidation resistance of  the  alloys could be improved because of the existence of  the three-phase alloys [6]. Huiren Jiang found that  with the addition of Ti content increased in the alloy,  the high-temperature oxidation  resistance of Nb based alloy became better, and the oxidation products  in Nb-based high-temperature alloy were TiO2, Nb2  O5, Ti2Nb10O29 and TiNb2O7. The PBR (Pilling-Bed  worth Ratio) represents the ratio of elementary cell  volume of metal oxide to the elementary cell volume of   the equivalent metal where the oxide has been created.  The increase of Ti content reduced the PBR value of  the alloy oxide and could improve the integrity of  oxide film [7]. ZrB2 is attracting more and more attention due to  its high melting point (3027°C), low theoretical density (6.1 mg/m3), high hardness (23 Gpa), good thermal  conductivity  (60-120 W/mK)  and  electrical  conductivity (~107), excellent physical and chemical  stability under high  temperature, neutron capture,  corrosion resistance, light weight, etc. [8,9]. The excellent high-temperature performance makes  it commonly  to be used  in  the ultra-high  temperature  materials. Although ZrB2 has high melting point, its  oxidation  resistance  is  poor  over  1000°C.  Thermogravimetry  test proves  that  there will be  a loss of weight when the ZrB2 is heated to 1000°C in  the air, the reaction of ZrB2 with oxygen generates  ZrO2 and volatile B2O3, the remaining ZrB2 on the  surface is not enough to form a dense oxide film and it  will continue to be oxidised. The researchers usually  solve this problem by adding a second phase into ZrB2  such as SiC [10]. In previous studies, it was reported that the compressive strength of the composites decreases with the  increase of the experimental temperature for NbMo ZrB2 composites. When the temperature is constant,  the compressive strength of the composites increases  with the increase of ZrB2 content [11]. The elastic  modulus of composites is positively related to ZrB2  content, and fracture toughness and ZrB2 content are  negatively correlated [12]. The NbMo series materials   CONTACT Zongde Liu  lzd@ncepu.edu.cn  Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of  Education, North China Electric Power University, Beijing 102206, China  © 2020 Informa UK Limited, trading as Taylor & Francis Group  \\x0c', '74  Z. LIU ET AL.  have higher brittleness and less toughness, while Ti  can effectively improve the toughness of the materials.  In spite of all this work on the resistance of oxidation  and the mechanical behaviour of Nb-Ti alloy or composites with ZrB2,  the high-temperature oxidation  resistance of Nb matrix composites reinforced by  ZrB2 particles is far from being understood [2,13-16]. In order to discuss the effect of boride particles on  the oxidation resistance of Nb matrix materials, metal  matrix ceramic composites were made from Nb-Ti  alloy containing 20 wt.% Ti with different content of  ZrB2  (0-60 wt.%). The oxidation mass  increment,  XRD, SEM and EDS methods are used to investigate  the influence of ZrB2 content on the oxidation property of the composites.  Materials and methods  Experimental materials  The ceramic metal composites used in the experiment  were made by hot pressing sintering process. The  sintering temperature was 1600°C and the pressure  was 30 Mpa. The specimens composition are presented in Table 1. The specimens were cut into a size  of 10 mm×8 mm×5 mm. Z0 is Nb-Ti alloy with 20 wt.  % Ti which is considered to be a comparative specimen. All the specimens were cleaned for 30 minutes  with alcohol in the ultrasonic cleaner and then dried.  Experimental method  High-temperature oxidation experiments were conducted in the model SK-6-12 tube electric resistance  furnace (6 kW)  in air at 800°C, 1000°C, 1200°C,  respectively. The specimens were placed in an alumina  crucible  so  that  the  total mass change  including  spalled  scales  could also be obtained. The mass  changes of the specimens with possible spalled scales  contained  in the crucible after oxidation test were  weighed in an electronic balance with an accuracy of  0.0001 g. The oxidation time of the specimens as  shown in Table 2. The oxidation weight gain of every  specimens was measured three times at each time  point, and the average was taken. Kinetic curves of  mass gain (G) and oxidation time (t) were made. The mass gain G is calculated as shown in equation,   Table 1. Nominal compositions of NbTi matrix composites.  Mass Fraction/%  Specimen No.  Z0 Z1 Z2 Z3 Z4  Nb-Ti (wt. %)  100(20%Ti+80%Nb) 85(20%Ti+80%Nb) 70(20%Ti+80%Nb) 55(20%Ti+80%Nb) 40(20%Ti+80%Nb)  ZrB2 (wt. %)  0 15 30 45 60  Table 2. The oxidation time of specimens at 800°C, 1000°C,  1200°C.  Specimen No.  Z0 Z1 Z2 Z3 Z4  800°C  28 h 96 h 96 h 96 h 96 h  Oxidation Time  1000°C  1200°C  28 h 96 h 96 h 96 h 96 h  6 h 60 h 60 h 60 h 60 h  G ¼ Δm=S  (1)   where G is the mass gain, g/mm2; is the total mass of  oxidised specimen increment, g; S is the total surface  area of the specimen, mm2. The oxidised specimens were analysed by X-ray  diffraction (XRD, Japanese Science D/max) and scanning electron microscopy (SEM, Quanta 200 F field  emission and SU8010 cold field emission) equipped  with an energy dispersive spectroscopy (EDS).  Results and discussion  Oxidation kinetics of the specimens  Figure 1 is oxidation kinetics curves of Z0, Z1, Z2, Z3,  Z4 at 800°C (a), 1000°C (b), 1200°C (c) in air. Table 3  is ultimate oxidation weight gain and Table 4 is average oxidation rate of NbTi matrix composites with  different ZrB2 contents at 800°C, 1000°C, 1200°C,  respectively. It can be seen that the oxidation mass  gain of per unit area increases with time. As time  increases, the oxidation weight gain of every specimens gradually changes from fast to gentle. And with  the rising of oxidation temperature, the weight gain  obviously increases. Figure 1 and Table 3 suggest that after oxidation at  800°C for 96 h, every specimens maintain relatively  mild mass gain. From Table 4  it  is clear that the  oxidation resistance order is Z4 (with 60% ZrB2) >Z3  (with 45% ZrB2) >Z1 (with 15% ZrB2) >Z2 (with 30%  ZrB2) >Z0(with 0% ZrB2) at 800°C. It is worth noting  that Z2 shows lower oxidation resistance than Z1 at  800°C. When the temperature rises to 1000°C, as we  can see from Figure 1, Tables 3 and 4 that the oxidation resistance of specimens increase with the increase  of ZrB2 content. Thus, the order of oxidation resistance is Z4 (with 60% ZrB2) >Z3 (with 45% ZrB2) >Z2  (with 30% ZrB2) >Z1 (with 15% ZrB2) >Z0(with 0%  ZrB2). At 1200°C, it can be seen from Figure 2 that  about the first 15 h of the oxidation process,  the  increase of ZrB2 content enhances the oxidation resistance of the specimens. But 17 h later, the oxidation  weight gain of Z3 has a sharp increase than Z2. When  oxidised for 40 h, the oxidation mass gain of Z4 has  a similar sharp increase and Z4 exceeds that of Z2 in  the end. The final oxidation weight gain at 1200°C is   \\x0c', 'MATERIALS AT HIGH TEMPERATURES  75  Figure 1. Oxidation kinetics curves of Z0, Z1, Z2, Z3, Z4 at 800 °C(a), 1000 °C(b), 1200 °C(c) in air.  Table 3. The ultimate oxidation weight gain value G (g/m2) of  T-Z composites with different ZrB2 contents at 800°C, 1000°C,  1200°C.  Temperature  (°C)  800 1000 1200  Sample No.  Z0  121.13 1337.13 1455.14  Z1  290.85 2768.08 2641.61  Z2  387.44 666.38 1468.74  Z3  159.86 327.32 2062.41  Z4  109.62 262.43 2202.38  Table 4. Oxidation rate (K/g/(m2·h)) of T-Z composites with  different ZrB2 contents at 800°C, 1000°C, 1200°C.  Temperature  (°C)  800 1000 1200  Sample No.  Z0  4.33 44.57 242.52  Z1  3.03 28.83 83.60  Z2  4.04 6.94 45.90  Z3  1.67 3.41 64.45  Z4  1.14 2.73 68.82  ordered as Z2 (with 30% ZrB2) >Z3(with 45% ZrB2)  >Z4 (with 60% ZrB2) >Z1 (with 15% ZrB2) >Z0(with  0% ZrB2).  XRD and surface oxidation morphology analysis  The oxidation products of the five specimens analysed  by XRD at 800°C, 1000°C and 1200°C are displayed in  Table 5. Figure 2 shows the surface oxide XRD patterns of  Z0 at 800°C (a), 1000°C (b) and 1200°C (c). It can be  seen that the oxidation products of Z0 after oxidation  at 800°C are Ti2Nb10O29, Nb2O5, TiO2 and TiNb2O7.  The products at 1000°C and 1200°C are Nb2O5 and  TiNb2O7. Nb2O5  is not dense which has poor   Figure 2. Surface oxide XRD patterns of Z0 at 800 °C(a),1000 °  C(b) and 1200 °C(c).  Table 5. The oxidation products of specimens at 800°C, 1000°C,  1200°C.  Sample  No.  Z0  Z1  Z2  Z3  Z4  800  TiNb2O7, Nb2O5, Ti2  Nb10O2, TiO2 TiNb2O7, Nb2O5, Ti2  Nb10O29, TiO2 Nb2O5, Ti2Nb10O29,   TiO2 Nb2O5, TiO2  Nb2Zr6O17  Nb2O5, TiO2  Nb2Zr6O17  Temperature/°C  1000  TiNb2O7, Nb2  O5 TiNb2O7, Nb2  O5, TiO2 TiNb2O7, Nb2  O5, TiO2 TiNb2O7, Nb2  O5,TiO2,  Nb2Zr6O17 Nb2O5, TiO2  Nb2Zr6O17  1200  TiNb2O7,Nb2O5  TiNb2O7, Nb2O5, TiO2  TiNb2O7, Nb2O5, TiO2   Nb2Zr6O17 TiNb2O7, Nb2O5, TiO2,   Nb2Zr6O17  TiNb2O7, Nb2O5, TiO2  \\x0c', '76  Z. LIU ET AL.  oxidation resistance. The reaction between Nb2O5 and  TiO2 produces stable and compact structure which is  TiNb2O7  [17]. But  the structure of Ti2Nb10O29  is  looser than TiNb2O7. Shown in Figure 3 is the XRD patterns which is  conducted on the surfaces of the scales formed on the  specimens Z1, Z2, Z3, Z4 at 800°C, 1000°C,1200 °C for  96 h, 96 h, 60 h, respectively. It can be seen that the  oxidation products of Z1  (Figure 3(a)) are mainly   TiNb2O7, Ti2Nb10O29, Nb2O5, TiO2. Compared  to  Z1,  the oxidation products of Z2 do not contain  TiNb2O7. It may be the reason that the oxidation  resistance of Z2 is worse than Z1. And the oxidation  products of Z3 and Z4 are Nb2Zr6O17, TiO2 and Nb2O5  at 800°C. While ZrO2 can form the network skeleton  structure of the material and react with Nb2O5 to form  a thin film-like structure which is Nb2Zr6O17. And the  oxygen diffusion resistance increases, so that the oxide  film can still be more closely integrated. The oxidation products of specimen Z1 after oxidation at 1000°C for 96 h (Figure 3(b)) are mainly Nb2O5  and TiNb2O7, with a very limited amount of TiO2. As  the ZrB2 content increases to 45%, the oxide content of  Nb and Zr increased. They are mainly Nb2Zr6O17 and  Nb2O5, a small amount of TiNb2O7 and TiO2. Several  oxides form a mixed oxide scale, which can improve  the adhesion and densification of the oxide scale and  prevent oxygen diffusion. When the oxidation temperature rises to 1200°C  (Figure 3(c)), the oxidation products of Z1 after oxidation for 60 h are TiNb2O7, Nb2O5 and TiO2. The  products of Z2 are TiNb2O7, Nb2O5, TiO2 and Nb2  Zr6O17. And for Z3, they are TiNb2O7, Nb2Zr6O17,  TiO2 and Nb2O5. Specimen Z4 are mainly TiNb2O7,  Nb2O5 and TiO2, a very little Nb2Zr6O17. The XRD results indicate that with the increase of  ZrB2 content, the temperature of the formation of  TiNb2O7 becomes higher. A possible explanation is  that when ZrB2 content rises, the diffusion of oxygen  atoms is hindered. So the oxidation resistance of the  specimens increases.  The microstructure analysis of oxide scale  The secondary electrons and the characteristic X-ray  signals of the oxidised specimens are collected to analyse the surface topography and the micro-domain  composition of the specimens.  Figure 3. Surface oxide XRD patterns of T-Z composites after  96 h oxidation at 800 °C(a), 96 h oxidation at 1000 °C(b) and  60 h oxidation at 1200 °C(c).  Figure 4. Macro morphology graph of Z0 after oxidation for  28 h at 1000 °C.  \\x0c', 'MATERIALS AT HIGH TEMPERATURES  77  The macro morphology graph of Z0 after oxidation  for 28 h at 1000°C is displayed in Figure 4, as we can  see in the figure that a large number of oxide scale  spalls. The higher the temperature, the surface of the  oxide layer shells off more in the cooling process. The  morphologies of the surface oxide layer are displayed  in Figure 5(a,b), it can be seen that there are many  holes and cracks on the surface and the oxide film is  not compact. As we know from the XRD results, the  oxide scales consist of mixed oxides. The mixed oxides   cannot be determined at present because these particles are very small. When ZrB2 is added in the composites, the microstructure of oxidised specimens Z1, Z2, Z3, Z4 are compared. Figure 6  shows  the  surface morphology of  specimens Z1, Z2, Z3, Z4 at 800°C for 96 h. The scales of  Z1 and Z2 (Figure 6(a,b)) have more surface cracks on  the oxide film. But Z3 and Z4 (Figure 6(c,d)) are relatively dense, and the scales are homogeneous and compact. At 1000°C, with the increase of ZrB2 content, as we   Figure 5. SEM images of surface morphologies of Z0 at different magnifications after oxidation for 28 h at 1000 °C.  Figure 6. SEM images of the surface morphology of specimens Z1, Z2, Z3, Z4 after oxidation for 96 h at 800 °C.  \\x0c', '78  Z. LIU ET AL.  can see in Figure 7, the oxide film on the surface of the  specimens gradually becomes dense. The surface topography changes from loose to complete. The oxidation  resistance of the materials is gradually increased. The  SEM surface topography displays in Figure 8 and the  EDS analysis results of the chemical composition of the  selected points have been presented in Table 6. It can be  seen in Figure 7(a) that the scale formed on the specimen Z1 is loose and characterised by a large number of  irregularly shaped small particles of 1-2 μm in size. And  there are large gaps between the particles. Oxygen ions  invade the interior of the alloy through the pore gap at  high temperature to accelerate the oxidation, and the  oxide film does not have a good protective effect. EDS  analysis (Table 6) reveals that is TiNb2O7 (point A). As  the ZrB2 content increases to 30 wt.% in Figure 7(b), the  scale surface consists of two phases. The diameter of  dark phases with irregular shape particles (point B) is  2-3 μm which is mainly Nb and Ti mixed oxide. And  the particle gap is smaller than Z1. Point C distributes  very fine particles which are light-grey and there is  almost no gaps between the particles. It can be speculated from the EDS that the light-grey phases distributing in the dark matrixes are mainly Nb2Zr6O17. As  shown in Figure 7(c) which contains 45 wt.% ZrB2, it   is obvious that the oxide film is formed by the massive  particles and the rod-shaped grains. The rod-like crystals (point D) with a length of about 10-20 μm are Nb2  O5. The smaller white bright particles (point E) are  TiNb2O7 and the particles in point F are Nb2Zr6O17.  This  is  the  same  result  as XRD. Figure 7(d)  is  a specimen containing 60 wt.% ZrB2. The scale consists  of grey irregular phase and light-grey fine particles on  the surface. It can be seen from EDS that the agglomerate particles (point H) are Nb2O5 and the fine particles (point G) are Nb2Zr6O17. And this is consistent  with the results obtained by XRD. When the oxidation temperature rises to 1200°C, as  shown in Figure 8, the scale of Z1 after oxidation for  32 h has a large number of cracks and holes but Z2, Z3,  Z4 are relatively dense. It can be seen that there are two  kinds of phases which are bulk phases and lamellar  phases in specimens Z3 and Z4. Table 7 is the EDS  analysis results of the chemical composition of the  selected points. It can be speculated from the EDS  that lamellar phases are mainly Nb2Zr6O17 (Figure 9  (e)) and the bulk phases are TiNb2O7 (Figure 9(f)).  Earlier reports show that when the temperature  is  higher than 1000°C, a large number of ZrB2 react as  follows.   Figure 7. SEM images of the surface morphology of specimens Z1(a), Z2(b), Z3(c), Z4(d) after oxidation for 96 h at 1000 °C.  \\x0c', 'MATERIALS AT HIGH TEMPERATURES  79  Figure 8. SEM images of the surface morphology of specimens Z1(a), Z2(b), Z3(c), Z4(d)after oxidation for 32 h at 1200 °C.  Table 6. Chemical composition of the sites marked with letters  in Figure 8, determined by EDS analysis.  Mass fraction/%  Position  A B C D E F G H  O  16.05 16.11 14.63 16.78 24.02 21.48 10.96 14.08  Zr  05.71 10.02 33.22 03.92 13.27 40.78 54.31 05.94  Nb  66.75 38.54 37.44 69.14 29.15 26.94 30.05 69.47  Ti  11.50 35.33 14.71 10.16 33.57 10.81 04.68 10.50  Table 7. Chemical composition of the sites marked with letters  in Figure 9, determined by EDS analysis.  Mass fraction/%  Position  A B C D E F G  O  13.57 20.38 17.97 30.30 22.11 13.10 28.53  Zr  10.39 10.80 07.13 38.54 57.56 15.30 54.50  Nb  60.07 56.26 63.80 14.54 17.17 13.92 14.48  ZrB2 þ 5O2 ðg Þ ! ZrO2 þ B2O3  B2O3 ðI Þ ! B2O3 ðg Þ  Ti  15.97 12.55 11.09 16.62 03.16 57.68 02.05  (2)   (3)   They are oxidised to form ZrO2 and liquid B2O3. As  the oxidation reaction to generate liquid B2O3 rate is  less than the rate of generating gaseous B2O3, the  liquid B2O3 evaporates, a continuous liquid B2O3 protective layer cannot be formed and the matrix cannot  be protected [18]. The large volume expansion and the  loss of liquid phase caused by the conversion of ZrB2  to ZrO2 creates cracks and pores which exacerbate  oxidation. Therefore, Z3 and Z4 with higher ZrB2  contents show poorer oxidisability than Z2. Z1 has  large  internal stress of oxide  layer because of  the  high Nb content. The oxide layer is easy to fall off  and so that the protective effect of TiNb2O7 is weak.  Thus, the oxidation resistance of Z1 at 1200°C is poor.  Conclusions  ZrB2-NbTi composites were hot-pressed at 1600°C  under the pressure of 30 Mpa with ZrB2 content ranging from 0 to 60 wt.%. The object is to study their  oxidation characteristics.  (1) Compared with the Nb-Ti alloys, the addition  of ZrB2 enhances the oxidation resistance of it.   \\x0c', '80  Z. LIU ET AL.  Figure 9. SEM images of the surface morphology of specimens Z1(a), Z2(b), Z3(c), Z4(d) after oxidation for 32 h at 1200 °C.  After ZrB2  is oxidised, ZrO2 and B2O3 are  formed into a networked skeleton of the material so that the scale is more tightly bound to the  substrate. The temperature required to form  TiNb2O7  is higher as  the ceramic content  increases. At the same oxidation temperature,  the order of the protective effect of the oxidation products on  the matrix  is Nb2Zr6O17  > TiNb2O7 > Ti2Nb10O29 > TiO2 > Nb2O5. (2) The order of oxidation resistance is Z4 (with 60%  ZrB2)>Z3  (with  45%  ZrB2)>Z1  (with  15%  ZrB2)>Z2 (with 30% ZrB2) >Z0 (with 0% ZrB2)  at 800°C. Compared to the oxidation products of  Z2 with 30% ZrB2, the specimen Z1 contains  TiNb2O7 which has a compact structure that it  can protect the oxide film well. Thus, Z1 shows  better oxidation resistance to Z2. When the ZrB2  content continues to rise to 45 wt.% and 60 wt.%,  due to the reaction of ZrO2 with Nb2O5 to form  Nb2Zr6O17 which has a good protection for oxide  film, the oxygen diffusion resistance increases and  the oxide film is not easy to rupture and spall. The  integrity of the oxide film improves.  (3) At 1000°C, the order of oxidation resistance is Z4  (with 60% ZrB2)>Z3 (with 45% ZrB2)>Z2 (with  30% ZrB2)>Z1 (with 15% ZrB2) >Z0 (with 0%  ZrB2). Nb2Zr6O17 plays a major protective role  in the oxide film with the increase of ZrB2 content  and the oxidation resistance of NbTi-matrix composites gradually increases. (4) When the temperature reaches 1200°C, the order  of oxidation resistance is Z2 (with 30% ZrB2)>Z4  (with 45% ZrB2)>Z3 (with 60% ZrB2)>Z1 (with  15% ZrB2) >Z0 (with 0% ZrB2). The higher the  content of ZrB2, the more ZrO2 and B2O3 are  oxidised. Due to the vigorous volatilisation of B2  O3 and the volume effect of ZrO2, it causes a large  number of cracks and holes which aggravate the  oxidation of the matrix. The oxide layer of Z1 with  15% ZrB2 is easy to fall off due to the large internal  stress which leads to poor oxidation resistance at  1200°C.  Disclosure statement  No potential conflict of interest was reported by the authors.  \\x0c', 'Funding  This work was supported by the National Natural Science  Foundation of China [11372110].  ORCID  Zongde Liu  Yuan Gao   http://orcid.org/0000-0001-6082-4587 http://orcid.org/0000-0003-0590-4802  References  [1] Zhao L, Guo X, Jiang Y. 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Influence of addition on  high temperature oxidation resistance of Nb-based  alloys. Chin J Nonferrous Met. 2014;20(8):2044-2049.  MATERIALS AT HIGH TEMPERATURES  81  [8] Song J, Du D, Xu M, et al. Oxidation behaviour of  ZrB2  matrix  composites  materials  at  high-temperature  conditions.  Powder  Metall  Technol. 2015;33(5):336-340,364. [9] Cao X. Synthesis of ZrB2 powder and preparation of  ZrO2 matrix composites reinforced by ZrB2. Beijing  University of Chemical Technology; 2012. [10] Yang W. Research on the growth and reinforced  mechanisms of  in situ TiB whiskers  in ZrB2-SiC  ceramic  joints. Harbin  Institute of Technology;  2014. [11] Wang Q, Liu Z, Hu H, et al. Experimental study on  preparation and properties of zirconium boride reinforced Nb Mo-matrix composites. Rare Metal Mater  Eng. 2017;46:2433-2436. [12] Liu Z, Wang Q, Gao Y, et al. Preparation and properties of hot-pressed NbMo-matrix composites reinforced with ZrB2 particles. 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},{
  "_id": 253,
  "PDF": "Synthesis and oxidation of nanocrystalline HfB2.pdf",
  "Text": "['Journal of Alloys and Compounds 368 (2004) 353-356  Synthesis and oxidation of nanocrystalline HfB2  Luyang Chen a , Yunle Gu a , Liang Shi a , Zeheng Yang a , Jianhua Ma a , Yitai Qian a,b,∗  a Department of Chemistry, University of Science and Technology of China, Hefei, Anhui 230026, PR China b Structure Research Laboratory, University of Science and Technology of China, Hefei, Anhui 230026, PR China  Received 17 July 2003; received in revised form 25 August 2003; accepted 25 August 2003  Abstract  Nanocrystalline hafnium diboride (HfB2 ) has been prepared by the reaction of HfCl4 with NaBH4 at 600 C in an autoclave. The X-ray powder diffraction (XRD) pattern can be indexed as hexagonal HfB2 with the lattice constants of a = 3.146 and c = 3.456 Å. The transmission electron microscopy (TEM) image shows a particle morphology with an average size of 25 nm. The selected area electron diffraction (SAED) pattern con firms the presence of hexagonal HfB 2 . The oxidation behavior of HfB2 is studied by thermogravimetric analysis (TGA) and differential thermal analysis (DTA). © 2003 Elsevier B.V. All rights reserved.     Keywords: Intermetallics; Chemical synthesis; X-ray diffraction  1.  Introduction  Transition metal borides are unique materials which have high melting points, high hardness and also high electric conductivity [1,2], and they are expected to be useful for a variety of ceramics and metals, coating materials, electron emitters and catalysts. Hafnium diboride (HfB2 ) has an AlB2 -type hexagonal structure, which can be used in thin film resistors [3] and explored as diffusion barriers in microelectronics [4-7] . Recently, HfB2 plus silicon carbide offers a good combination of properties that make it a candidate for application in airframe leading edges on sharp-bodied reentry vehicles [8]. Hafnium boride powder can be synthesized by various methods. Reich et al. prepared hafnium boride thin-films by plasma enhanced chemical vapor deposition [9]. Mikami et al. synthesized HfB2 by a reactive ion plating method, which used B2H6 -He mixture gas and metallic Hf as reactants [10]. Munir described the reaction processes of hafnium boride by the self-propagating high-temperature synthesis (SHS) [11]. However, it is difficult to obtain HfB 2 nanomaterials by these methods. Herein, we report a convenient preparation method of nanocrystalline HfB2 by the reaction of HfCl4 and NaBH4  ∗  Corresponding author. Tel.: +86-551-3601589; fax: +86-551-3607402. E-mail address: chendavy@mail.ustc.edu.cn (Y. Qian).  0925-8388/$ - see front matter  © 2003 Elsevier B.V. All  rights reserved.  doi:10.1016/j.jallcom.2003.08.086     at 600 C in an autoclave, used as source materials.  in which HfCl4 and NaBH4 are  2. Experimental procedure        Firstly, excessive analytical grade NaBH4 powders (30 mmol) and liquid analytical pure HfCl4 (5 mmol) were placed into a stainless steel autoclave with a quartz tube. And then, the autoclave was sealed under Ar atmosphere and heated at 600 C for 12 h, followed by cooling to room temperature. The product was washed with distilled water and absolute ethanol for several times to remove the impurities. The final product was vacuum-dried at 60 C for 4 h. X-ray powder diffraction (XRD) pattern was carried out on a Rigaku Dmax-␥A X-ray diffractometer with Cu K␣ radiation (λ = 1.54178 Å). The morphology of nanocrystalline HfB2 was examined from transmission electron microscopy (TEM) images taken with a Hitachi H-800 transmission electron microscope. X-ray photoelectron spectra (XPS) were recorded on a VGESCALAB MKII X-ray photoelectron spectrometer, using non-monochromatized Mg K␣ X-rays as the excitation source. Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) profiles were collected with a Shimadzu-50 thermoanalyzer apparatus under flowing air.  \\x0c', '354  L. Chen et al. / Journal of Alloys and Compounds 368 (2004) 353-356  )  0 0 1  (  )  1 0 1  (  )  1 0 0  (  .  u  .  a  /  y  t  i  s  n  e  t  n  I  )  0 1 1  (  ) )  2 1 0 1 1 1  ( (  )  0 0 2  (  )  1 0 2  (  )  2 0 0  (  20  30  40  50  60  70  80  2Theta/degrees  Fig. 1. XRD pattern of  the HfB2  sample prepared by reaction of HfCl4 and NaBH4  for 12 h at 600     C.  3. Results and discussion  Fig. 1 shows the XRD pattern of the product. In Fig. 1, all of the nine peaks at d-spacings of 3.457, 2.725, 2.137, 1.728, 1.575, 1.464, 1.430, 1.364 and 1.266 Å can be indexed as hexagonal HfB2 ((0 0 1), (1 0 0), (1 0 1), (0 0 2), are a = 3.146 Å and (1 1 0), (1 0 2), (1 1 1), (2 0 0), (2 0 1)). The lattice constants c = 3.456 Å, in good agreement with a = 3.140 Å and c = 3.470 Å (JCPDS card #75-1049). No evidence of Hf, HfO2 and other impurities are observed. The broadening of the diffraction peaks indicates that the sample is nanoscale in size. The crystallite size is estimated to be about 25 nm in diameter, when using the Debye-Scherrer equation with a shape factor of 0.9.  The transmission electron microscope (TEM) image and a selected area electron diffraction (SAED) pattern are shown in Fig. 2. As shown in Fig. 2a, the product exhibits small particle morphology. The particle size is estimated in the range of 20 -30 nm in diameter, which is in agreement with the value calculated by the Debye-Scherrer equation. In Fig. 2b, the diffraction rings from the inner to the outer region occur, at d-spacings of 3.46, 2.73 and 2.14 Å, matching HfB 2 (0 0 1), (1 0 0), (1 0 1) planes, in good agreement with the XRD results. To give further evidence for the formation of HfB2 , the binding energy of the electrons at the Hf 4f5/2 level was investigated. The peak at 17.70 eV in the XPS spectra (as shown in Fig. 3) corresponds to the Hf-B bond in HfB2 , in  Fig. 2. TEM image and ED pattern of  the HfB2  sample prepared by reaction of HfCl4 and NaBH4  for 12 h at 600     C.  \\x0c', 'L. Chen et al. / Journal of Alloys and Compounds 368 (2004) 353-356  355  Hf4f5/2  .  u  .  a  /  y  t  i  s  n  e  t  n  I  12  14  16  18  20  22  24  Binding Energy/eV  Fig. 3. XPS spectra of  the HfB2  sample prepared by reaction of HfCl4 and NaBH4  for 12 h at 600     C.  accord with the data in the literature (Hf 4f5/2 : 16.40 eV) [12]. No obvious peaks ascribable to the Hf-O bond were observed. The TGA-DTA technique was used to analyze the thermal behavior of the HfB2 sample. A typical TGA-DTA profile studied in flowing air is shown in Fig. 4. During the initial heating between 27 and about 100 C, water is removed from the sample corresponding to an initial weight loss of 3.51%. There is a weight gain starting at about 400 C. From 400 to 1000 C the weight of the sample increases by about                 22% and a corresponding exothermic peak is observed near 723 C. This weight gain can be ascribed to HfB2 oxidation to form HfO2 and B2O3 [13]. The weight gain suddenly becomes slow at about 770 C. Around 1000 C B2O3 evaporates extensively [14], which counteracts the weight gain due to the oxidation of HfB2 , so the weight tends to keep in balance. This B2O3 evaporation results in a final weight gain (122%) that is less than the theoretic one (140%). The possible formation mechanism of HfB2 has been proposed. In the process to nanocrystalline HfB2 by reacting     120  115  %  /  n  i  a  G  t  h g  i  e  W  110  (b)  105  100  95  (a)  0  200  400  600  Temperature/Deg C  800  1000  Fig. 4. TGA (a), DTA (b) profile of  the HfB 2  sample prepared by reaction of HfCl4 and NaBH4  for 12 h at 600     C.    \\x0c', '356  L. Chen et al. / Journal of Alloys and Compounds 368 (2004) 353-356     HfCl4 with NaBH4 at 600 as follows: HfCl4 + 2NaBH4 → HfB2 + 2NaCl + 2HCl + 3H2  C, the reaction can be described  (1)     The NaBH4 began to decompose with increasing temperature (>500 C) [15] as expressed in Eq. (boiling point = 432 (2). The BH3 reacted further with HfCl4 C) to produce HfB2 and HCl, in which HCl immediately reacts with NaH to produce H2 and NaCl (Eq. (3)): NaBH4 → BH3 + NaH HfCl4 + 2NaH + 2BH3 → HfB2 + 2HCl + 3H2  (2)  (3)           The reaction temperature and time play very important roles in the formation of hexagonal HfB2 . When the temperature was higher than 500 C, the crystallinity and the crystallite size both increased. HfB2 did not form when the temperature was lower than 450 C. The reaction was usually incomplete and the crystallinity was poor when the reaction time was less than 6 h. However, varying the reaction time between 12 and 24 h did not significantly affect the crystallinity and the particles size. The maximal pressure is about 20 MPa at 600 C, which is estimated according to the amount of hydrogen treating as ideal gases. It is believed that the increasing pressure in the autoclave may be beneficial for the formation of HfB2 .     4. Conclusions  In summary, nanocrystalline HfB2 particles of 25 nm size with the hexagonal structure have been successfully prepared by a convenient reaction of HfCl4 with NaBH4 at 600 C for 12 h. Our study demonstrates that a higher reaction temperature is beneficial for the crystallinity and for an     increased crystallite size of the samples. The oxidation of nanocrystalline HfB2 in flowing air begins to proceed at a temperature of 400 C.     Acknowledgements  This work is supported by the Chinese National Science Research Foundation.  References  [1] R.M. Adams, Boron, Metallo-Boron Compounds  and Boranes,  In terscience, New York, 1964.  [2] G.V. Samsonov, I.M. Vinitskii, Handbook of Refractory Compounds,  Plenum Press, New York, 1980.  [3] D.S. Wuu, M.L. Lee, T.Y. Lin, R.H. Horng, Mater. Chem. Phys. 45  (1996) 163.  [4] J.R. Shappirio,  J.J. Finnegan, R.A. Lux,  J. Vac. Sci. Technol. B4  (1986) 1409.  [5] C.S. Choi, G.A. Ruggles, A.S. Shah, G.C. Xing, C.M. Osburn, J.D.  Hunn, J. Electrochem. Soc. 138 (1991) 3062.  [6] G. Sade, J. Pelleg, Appl. Surf. Sci. 91 (1995) 263.  [7] E. Kolawa,  J.M. Molarius, W. Flick, C.W. Nieh, L. Tran, M.A.  Nicolet, F.C.T. So, J.C.S. Wei, Thin Solid Films 166 (1988) 29.  [8] S.R. Levine, E.J. Opila, M.C. Halbig,  J.D. Kiser, M. Singh,  J.A.  Salem, J. Eur. Ceram. Soc. 22 (2002) 2757.  [9] S. Reich, H. Suhr, K. Hanko, L. Szepes, Adv. Mater. 4 (10)  (1992)  650.  [10] H. Mikami, S. Takahashi, T. Sato, K. Shimakage, Denki Kagaku  62 (8)  (1994) 686.  [11] Z.A. Munir, Metall. Trans. A 23 (1992) 7.  [12] C.L. Perkins, R. Singh, M. Trenary, T. Tanaka, Y. Paderno, Surf.  Sci. 470 (3)  (2001) 215.  [13] S. Takahashi, T. Sato, K. Shimakage, Denki Kagaku 63 (2)  (1995)  140.  [14] A. Tampieri, A. Bellosi, J. Mater. Sci. 28 (1993) 649.  [15] A.G. Ostroff, R.T. Sanderson, J.  Inorg. Nucl. Chem. 4 (1957) 230.  \\x0c']"
},{
  "_id": 254,
  "PDF": "Synthesis of ZrC–SiC Powders from Hybrid Liquid Precursors with Improved Oxidation Resistance.pdf",
  "Text": "['Synthesis of ZrC-SiC Powders from Hybrid Liquid Precursors with Improved  Oxidation Resistance  Xiao-Fei Wang,  ‡,§  Jia-Chen Liu,  §  Feng Hou,  §  Ji-Dong Hu,  ‡  Xin Sun,  ‡  and Yan-Chun Zhou  ‡,†,**  ‡  Science and Technology on Advanced Functional Composite Laboratory, Aerospace Research Institute of Materials &  Processing Technology, Beijing 100076, China  §  Key Laboratory of Advanced Ceramics and Machining Technology of Ministry of Education, School of Materials Science and  Engineering, Tianjin University, Tianjin 300072, China  ZrC-SiC powders are synthesized by high-temperature pyroly sis  of  hybrid  liquid  precursors,  which  are  prepared  from  organic Zr-containing precursor  (PZC) and liquid polycarbosi lane  (LPCS). Due  to  the  excellent miscibility  between PZC  and LPCS,  the hybrid liquid precursors are formed by dissolv ing PZC into LPCS without adding organic solvent. The  vis cosity  and  elemental  content  of Zr  and  Si  of  the  hybrid  precursors  are  readily  adjustable  by  controlling  the LPCS/  PZC mass ratio. SEM and TEM observations ZrC-SiC powders pyrolyzed at 1550°C exhibit  reveal  that  the  spherical mor phology with characteristic dimension of  less  than 60 nm, and  the two phases are uniformly distributed in composite powders. the ZrC-SiC powders  The  advantage  of  synthesized  by  this  novel method  is  demonstrated  by  investigating  the  oxidation  behavior of powders with different amount of SiC and ZrC. Below 700°C, ZrC quickly oxidizes to generate an almost nonprotective ZrO2 scale, whereas at ~ 1000°C, dense and protective SiO2 forms that improves the oxidation resistance of the ZrC-SiC composite powders.  I.  Introduction  U LTRAHIGH-TEMPERATURE  ceramics  (UHTCs), with  their  high melting points, good thermal stability, and chemi cal  inertness, have promising  applications as  leading edges, vehicles.1-4 continuous  nose  tips,  reentry  heat  shields  of  hypersonic  UHTCs  have  also  been  designed  as matrix  of  fiber-reinforced  ceramic  matrix  composites,  which  often  requires polymeric ceramic precursors in its polymer impregnation pyrolysis densification process.2,5-8 For ultrahigh-temperature applications, beside high melting point, good oxidation resistance is also a basic requirement.5,9 Zirconium  carbide  (ZrC)  exhibits  a unique  combination of properties,  such as high melting point, good thermal and electrical con ductivity, solid-state phase resistance.3,6,7 The  stability, and excellent  corrosion  rapid oxidation of ZrC under oxidizing  atmosphere, however, has hampered its widespread applica tions. Therefore,  improving the oxidation resistance of ZrC  has been the task of many works, which has led to significant  increase in papers focusing on adding SiC into ZrC, because  it  can significantly  improve  the high-temperature oxidation  resistance of ZrC by forming a protective amorphous silicate coating.10,11  The main methods for the synthesis of ZrC-SiC composite  powders are solid mixing, liquid precursor conversion, and polymer precursor derivation.12-15 Solid mixing is simple but making the composition and size homogeneity is the major  challenge. Liquid precursor conversion requires high pyrolytic  temperature, and overcome the structural and compositional  heterogeneity  is  also difficult. Polymer precursor derivation  has the following benefits: (1) It is easy to control the compo sition homogeneity in molecular level.  (2) It has high produc tivity with  a  relative  low fabricating  cost.  (3)  It  permits  homogeneous  phase  distribution without  the  necessary  for  adding extra solvent. (4) It has simple equipment requirements  and relatively low pyrolytic temperature.  (5)  It can be trans ferred to the substrate and/or coating for UHTCs-based com posites by adjusting the processing parameters.  In this work, we report a simple method for of ZrC-SiC composite powders  the synthesis  from hybrid liquid precur sors. The hybrid liquid precursors are prepared by dissolving  organic Zr-containing  precursor  (PZC)  directly  into  liquid  polycarbosilane  (LPCS) without  the necessary for  the addi tion of extra solvent. This new method is different from those  reported previously, wherein liquid PCS and PZC are used  as SiC and ZrC precursor,  respectively. Most  importantly,  unlike using conventional  solid PCS and PZC precursors,  in  this method no extra solvent  is used, because LPCS is a spe cial PCS that  in liquid state at  room temperature  (RT). To  demonstrate  the advantage of the method and the the synthesized ZrC-SiC powders,  superior  properties of  the miscibil ity behavior between PZC and LPCS is  studied first. Then,  the structural  features and viscosity of  the hybrid liquid pre cursors are characterized. Moreover,  the phase composition,  elemental composition, and microstructure of sized ZrC-SiC composite  the as-synthe powders  are  also  characterized.  Finally,  the oxidation behavior and the structure evolution the ZrC-SiC powders are investigated.  of oxide scales of  II.  Experimental Procedure  (1)  Synthesis of ZrC-SiC Powders  A vinyl-substituted LPCS, which is produced via Grignard  coupling reactions, and PZC are used as starting materials to synthesize ZrC-SiC composite powders. Compared to the nor mal  solid PCS,  the LPCS is a special PCS that  is yellow in  color in liquid state at RT and enrichment of carbon. The viscosity of LPCS at RT is 0.78 Pa\\x01s, The average molecular weight (Mn) and (Mw) of LPCS are 658 and 1449, respectively (X. F. Wang, J. C. Liu, F. Hou, X. P.  indicating excellent fluidity.  Lu, X.  Sun, Y. C. Zhou, unpublished data). The PZC is a linear Zr-O-Zr chain polymer with  [(C4H8O)Zr(acac)2]n acetylacetone (Hacac) as ligand. It is in solid state at RT (dark red in color) with molecular weight (Mw) of ~64870 and a softening point of 170°C.8 The oxygen in PZC can be removed by  carbothermal  reaction  during  high-temperature  pyrolysis  R. Koc—contributing editor  Manuscript No. 34410. Received January 16, 2014; revised September 10, 2014;  approved September 10, 2014.  **Fellow, The American Ceramic Society.  †  Author to whom correspondence should be addressed.e-mail: yczhou@imr.ac.cn  197  J. Am. Ceram. Soc., 98 [1] 197-204 (2015)  DOI: 10.1111/jace.13274  © 2014 The American Ceramic Society  Journal  \\x0c', 'because the LPCS is carbon rich. The hybrid liquid LPCS-PZC  precursors are prepared by dissolving solid PZC into LPCS,  and magnetically stirring (200 r/min) at RT into a homoge neous  liquid. During  stirring,  the  initial  suspension  blends  completely and become a bit  slimy sol with orange in color.  Therefore, LPCS-PZC precursors with  various LPCS/PZC  mass ratios, namely, 1.0, 1.5, 2.0, 2.5, 3.0, 3.5, and 4.0, are pre pared. Afterward, the thermal curing of hybrid LPCS-PZC is carried out at 170°C in an oven for 2 h under  precursors  argon protection, and the slimy sol transforms gradually into a  rubbery solid with pale brown in color. The subsequent high temperature pyrolysis is performed by heating the cured hybrid precursors at a rate of 2°C/min to the target 1550°C, argon atmosphere. Finally,  temperature of  and held at  that  temperature for 2 h in a flowing the ZrC-SiC composite powders  with black in color are obtained after  cooling down to RT  along with the furnace.  (2)  Characterization Methods  The structures of PZC, LPCS, and hybrid LPCS-PZC precur sors are analyzed by Fourier transform infrared spectroscopy  (FTIR; Nexu 6700, Nicolet, Madison, WI) with KBr disks for  solid samples and KBr plates  for  liquid samples. The LPCS,  cured LPCS, PZC, and hybrid LPCS-PZC precursors are sub jected to a thermal analyzer (NETZSCH STA 449 C, Selb, Ger5°C/min. The  many)  in flowing  argon at  a heating  rate of  viscosity of LPCS-PZC precursors is measured with a viscome ter (Brookfield DV-II, Middleboro, MA) at RT using a rotat ing velocity of 25 r/min.  A D/max-2400 X-ray diffractometer (Rigaku, Tokyo, Japan) with CuKa radiation at a scanning speed of 0.04°/step is used to determine phase composition and crystallinity of the as-synthesized ZrC-SiC powders. To quantitatively identhe ZrC-SiC powders, an  tify the  elemental  composition of  electron  probe  microanalyzer  (EPMA1610;  Shimadzu,  Tokyo,  Japan)  is  used. The morphological  observation  is  conducted on a scanning electron microscope (SEM; Apollo 300, CamScan, Cambridge, UK). Microstructure of ZrC-SiC 200 kV Tecnai G2F20  powders  is  also investigated using  a  high-resolution transmission electron microscope  (HRTEM;  FEI, Philips,  the Netherlands) by dispersing the powders  in  ethanol, dipping onto a copper grid and drying in air.  (3)  Oxidation Tests  The oxidation behavior of ZrC-SiC powders is performed in  a thermal analyzer  (NETZSCH STA 449 C) by heating the powder samples from ambient temperature to 1400°C at a rate of 5°C/min in a flowing air atmosphere. Isothermal oxi dation examinations are carried out 400°C-1300°C for  in the temperature range  of  2 h  in  air.  In  isothermal  oxidation  experiments,  the  specimens  are  protected  in  argon  before  reaching to the target  temperature, and then air is introduced  into the furnace by a computer-controlled switch from argon  to air once the target temperature is reached. The phase comthe oxidized ZrC-SiC powders  position and morphology of  are characterized by XRD and SEM, respectively.  III.  Results and Discussion  (1)  Characterization of Hybrid LPCS-PZC Precursors  (A)  Miscibility Behavior:  To investigate the miscibility  behavior between LPCS and PZC without using  extra  sol vent, a series of LPCS-PZC precursors are prepared by varying the LPCS/PZC mass ratio. The viscosity (g) of individual  LPCS-PZC precursor with each fixed mass  ratio is measured  at RT using a rotating viscometer under the same experimen tal  conditions. The  interval of  the  time between the  initial  suspension and the homogeneous solution is defined as prep aration time (t). As  listed in Table I, both the viscosity and  preparation  time  of  hybrid LPCS-PZC precursors  change  along a similar  trend. When the LPCS/PZC mass  ratio pro gressively increases from 1.0 to 4.0, to 1.82 Pa\\x01s, meanwhile, the viscosity significantly tion time markedly declines from 190 to 10 min. Our  decreases  from 19.06  the prepara study  indicates  that LPCS and PZC can be mutually soluble with out  the necessary  for  the  addition of  extra  solvent. More over,  the amount of LPCS is  found to be important  for  the  affecting  the  rheology  of  hybrid  LPCS-PZC  precursors  because LPCS is liquid state with excellent lower mass ratio of LPCS/PZC = 1.0,  fluidity. At  a  the hybrid LPCS-PZC  precursor shows relatively viscous, whereas at a higher mass ratio of LPCS/PZC = 4.0, displays excellent intermiscibility.  the hybrid LPCS-PZC precursor  Therefore,  the  viscosity  and  preparation  time  of  hybrid LPCS-PZC precursors  are  adjustable by varying the LPCS/PZC mass ratio.  (B)  Structural Characters  of Hybrid LPCS-PZC Pre cursors:  Figure 1(a) presents the functional groups of PZC. \\x001 is ascribed to the characterisThe broader band at 652 cm tic absorption band of Zr-O stretching vibration, suggesting are Zr-O-Zr that there PZC.6 The peaks at 933 and 1023 cm (C-CH3) + (C=O) and C-O stretching vibrations, \\x001 are assigned to respectively.7,17 The peaks at 1277 and 1595 cm the absorption bands of acetylacetone ligand.16 The peak at is attributed to C-H bending vibration, whereas \\x001 is due to stretching of C-H the hump observed at 2930 cm bond in -CH2, respectively.17 It is worth mentioning that the peak at 1533 cm is referred to C=C stretching from the enol form (C=C-OH) of acetylacetone, which is easy to transform into C-C=O and readily to chelate to zirconium, thus, a cyclic structure of the complex may form.6,7  segments  in the polymer  chain  of  \\x001  correspond to the  1384 cm  \\x001  \\x001  vibration  Figure 1(b)  illustrates  the  IR spectrum of \\x001 are assigned to LPCS. The  absorption bands at 765, 846, and 936 cm  Table I.  Effect of  the LPCS/PZC Mass Ratio on the  Viscosity and Preparation Time of Hybrid Liquid LPCS-PZC  Precursors  LPCS/PZC  mass ratio  liquid precursors, g (Pa\\x01s) Viscosity of hybrid  Preparation  time,  t  (min)  1.0  19.06  190  1.5  10.87  160  2.0  7.63  120  2.5  6.24  90  3.0  4.71  50  3.5  2.98  20  4.0  1.82  10  (a)  (b)  (c)  Fig. 1.  Fourier  transform infrared spectroscopy spectra of the (b) LPCS, and (c) hybrid LPCS/PZC = 1.0 in  precursors:  (a) PZC,  mass ratio.  198  Journal of  the American Ceramic Society—Wang et al.  Vol. 98, No. 1  \\x0c', 'Si-C stretching and Si-CH3 group.18 The Si-CH2-Si stretch\\x001, \\x001, ing at 1047 cm the C-H bending in Si-CH2 at 1356 cm \\x001) and the C-H stretching in Si-CH2 (2882, 2920 cm reveal existence of a Si-CH2-Si the chain as the backbone polymer precursor.19 There is Si-H. The which refers to characteristic and C-H stretching Si-CH3 \\x001, respectively. The absorption peak at 1630 cm CH=CH2 stretching) \\x001 suggest C-H 3078 cm the incorporation of vinyl groups into this LPCS precursor.20  in this  a  strong band at  2135 cm Si-CH3 1253 and  \\x001  bands  of  in  are  observed  at  2960 cm (C=C  \\x001  and  vibration  in  at  The structure of as-prepared hybrid LPCS-PZC precursor also identified by FTIR using LPCS/PZC = 1.0  is  in mass  ratio as a model precursor, as  shown in Fig. 1(c). One  can  see  from Fig. 1(c)  that  the  hybrid  LPCS-PZC precursor  exhibits both typical LPCS and PZC characteristics, which  demonstrates  that LPCS and PZC are  intermiscible without  using extra solvent  to form homogeneous  solution. Further more, a clear increase on the absorption peaks at 1533 and \\x001 can be observed in the hybrid LPCS-PZC precurconfirming that the prepared solutions are LPCS-PZC  1595 cm  sor,  homogeneous  phase  rather  than  a mixture  of LPCS  and  PZC.  (C)  Thermal Gravimetric Behavior:  To understand the  thermal behavior of the hybrid LPCS-PZC precursor, LPCS, 170°C LPCS cured at PZC = 1.0 in mass  in  argon,  PZC,  and  the  LPCS/  ratio are  subjected to TGA analysis, as  shown in Fig. 2(a). The LPCS loses weight of 4 wt% below 200°C, 200°C and 490°C, 18 wt% between above 490°C-900°C. The total weight 1200°C, which corresponds  and  14 wt%  loss  is 36 wt% below  to a  ceramic  yield of  64 wt%.  Compared with the uncured LPCS, the major mass loss of the LPCS cured in argon in the temperature range of 400°C- 600°C is  13 wt%,  owing  to  the  rearrangement  reactions  along with the volatile species like hydrogen (H2) and methane (CH4) evolution.21 The ceramic yield of the cured LPCS (82 wt%) is much higher than that of uncured LPCS (64 wt  %), which attributes  to the  existence of vinyl groups  in the  LPCS that  can be  cross-linked into infusible  structures dur ing pyrolysis and thus contributes  to the high ceramic yield.  The TGA curve of PZC displays  that a considerable weight in the 200°C-  loss (approximately 35 wt% loss) takes place 600°C region, which is mainly due organic side groups accompanied by the polymerization and  to the decomposition of  structural relaxation. The TGA curve of the LPCS-PZC preof LPCS/PZC = 1.0  cursor with mass ratio below 300°C, the evaporation of (2) between 300°C and 400°C, 400°C and 600°C, the removing of to the weight loss, and (4) above 600°C releases of methyl groups.22  presents  four  zones:  (1)  oligomers  are  dominated;  a  stable  area  is  observed;  (3)  between  vinyl groups  leads  the mass  loss  is  related to the  Notably,  it  can be  seen from Fig. 2(a)  that  the PZC has a  relatively low ceramic yield of about 37 wt%, whereas the hybrid LPCS-PZC precursor with LPCS/PZC = 1.0 reaches a high ceramic yield of about 69 wt%. As a result, the intro duction of  an equivalent mass of LPCS into the PZC can  obviously enhance the final ceramic yield.  To further  investigate the effect of LPCS/PZC mass  ratio  on the ceramic yield, TGA curves of  the hybrid LPCS-PZC  precursors with various LPCS/PZC mass ratios are compared  in Fig. 2(b). From Fig. 2(b),  the ceramic yield increases from  69 to 76 wt% when the LPCS/PZC mass ratio increases from 1.0 to 4.0. This can be understood that more Si-H bonds from LPCS react with the C=C bonds from PZC through the  hydrosilylation  process  under  thermal  cross-linking  at  ele vated temperatures, which contributes yield.23 From the above description,  to the  high ceramic  the hybrid LPCS-PZC  precursor  undergoes  cross-linking,  an  organic-to-inorganic  transition and the conversion to amorphous or even crystal line ZrC-SiC powders during pyrolysis  [Eq.  (1)], which can  be predicted as the following:  PZC (s) þ LPCS (l) ! ZrC-SiC (s) þ CO (g) þ CH4 (g)  (1)  (2)  Phase Composition, Elemental Composition, and  Microstructure of ZrC-SiC Powders  (A)  Phase Composition:  The XRD patterns  of ZrC SiC powders derived from various LPCS/PZC mass ratios after pyrolysis at 1550°C are shown in Fig. 3. When the mass ratio of LPCS/PZC is 1.0 [Fig. 3(a)], the XRD pattern shows that ZrC and b-SiC are  the predominant phase. Diffraction  peaks  corresponding to (111),  (200), (220), (311), and (222) 2h = 33.16°, 38.51°, 55.61°, planes of ZrC are 66.32°, 69.68°, and respectively. The three major peaks at 2h = 35.71° (111), 60.13° (220), and 71.92° (311), along with the weaker ones at 2h = 41.77° (200) and 75.18° (222), are 6H-SiC b-SiC, at in b-SiC19 or  identified at  attributed 2h = 33.87° (101) is assigned to stacking faults a-SiC. Besides ZrC and SiC,  to  and  the  shoulder  of  tetragonal zirconia (t-ZrO2) and (m-ZrO2) are also detected, indicating the presence of oxygen in the PZC [Figs. 1(a) and (d)] and  monoclinic  zirconia  free Hacac is volatilized due to tion.7,14 Graphite is not detected, implying that it is present in minute amount.8 Similarly,  the  saturation  of  chela either  in amorphous  form or  no diffraction peak from SiO2 SiO2 has been converted into SiC or remained in amorphous form.4 The decreased intensity of diffraction peaks of zirco is observed,  suggesting  that  nia (ZrO2) with [Figs. 3(b)-(d)] indicates  the  increased mass  ratio  of  LPCS/PZC  that ZrO2 uct and reacts with C from LPCS to form ZrC.  is an intermediate prod In addition,  (a)  (b)  Fig. 2.  Thermogravimetry (TG) curves of  the different  samples  (a),  and the hybrid LPCS-PZC precursors with various LPCS/PZC mass  ratios (b).  January 2015  Synthesis of ZrC-SiC Powders  199  \\x0c', 'with the  increasing of LPCS to PZC ratio,  the  intensity of  SiC peaks increases and that of ZrC decreases, and thus SiC  becomes the predominant phase.  (B)  Elemental Composition Determined  by EPMA:  The elemental composition of  the ZrC-SiC powders is deter mined  by  EPMA,  and  the  results  are  summarized  in  Table II. It  is obvious to see that  the elemental compositions  are basically consisting of Zr, Si, and C. The chemical composition of ZrC-SiC powders derived from the hybrid preLPCS/PZC = 1.0  cursor  with  in  mass  ratio  is  ZrSi3.345C4.139O0.520,  suggesting that  the C source  is  insuffi cient  to  entirely  reduce  ZrO2 compositions of samples derived from the hybrid precursor with the mass ratio of LPCS/PZC = 2.0, 3.0, and 4.0 are  [Fig. 3(a)].  The  chemical  ZrSi6.445C8.064O0.484,  ZrSi9.732C12.220O0.433,  and  ZrSi12.933  C16.727O0.359, respectively. The sufficient and even excessive C from higher amount of LPCS warrants the complete reduc tion of ZrO2. Thus, ZrO2 gradually transforms into ZrC [Figs. 3(b)-(d)], but a tiny amount of O can still be detected by EPMA. Previous works by Xiang et al.24 and Liu et al.25  demonstrate that  the O could be dissolved in ZrC lattice to  form a “zirconium oxycarbide solid solution”. In addition,  it  is noted that  the Si and C contents visibly increase with the  increasing of LPCS/PZC mass  ratio from 1.0 to 4.0, but  the  Zr content clearly decreases (Table II). Therefore, the element composition in ZrC-SiC powders is readily controllable  by varying the LPCS/PZC mass ratio.  Moreover,  residual  carbon  content  in  samples  derived  from different LPCS/PZC mass  ratios  is also given, as  listed  in Table II. At a low LPCS/PZC mass ratio, C is insufficient  to totally  reduce  the ZrO2 high LPCS/PZC mass ratio, because the LPCS is C rich,  in the  carbothermal process;  at  the  C from LPCS is  sufficient  to reduce ZrO2 mal process, and even excessive C is remained. As a result,  in the  carbother the carbon content  increases from 1.98 to 5.08 wt% with the  mass ratio of LPCS/PZC increases from 2.0 to 4.0.  (C)  Microstructure  of  ZrC-SiC  Powders:  Figure 4  shows the microstructure of ZrC-SiC powders derived from LPCS/PZC = 1.0 in mass ratio. SEM micrograph in Fig. 4(a) the ZrC-SiC powders exhibit spherical morphol reveals that  ogy with smooth surfaces. Although the agglomeration exists, a multitude of ZrC-SiC particles is dispersed relatively  regularly, with an average particle size of less than 60 nm. The crystallinity of nanoscale ZrC-SiC powders and the dis tribution  of ZrC and SiC phases can be seen images, as shown in Figs. 4(b)-(c). According to Fig. 4(b), the ZrC-SiC powders are nanocrystalline with dimensions of about 30-50 nm, which are consistent with the broad peaks  from TEM  in XRD patterns  [Fig. 3(a)]. Moreover,  the ZrC and  SiC  phases  evenly “mingle” each other,  indicating that  the ZrC  and SiC phases are uniformly distributed. Figure 4(c) displays the lattice fringes of ZrC and SiC nanocrystals in ZrC-  SiC powders. The  d-spacings  of  0.25 and 0.27 nm can the (111) plane of b-SiC (a = 0.436 nm) of ZrC (a = 0.470 nm), interface between ZrC and SiC is clean without amorphous  be  assigned  to  and  (111)  plane  respectively. And  the  phase.  The above results demonstrate that  the ZrC-SiC compos ite powders are successfully synthesized by high-temperature  pyrolysis of  the hybrid liquid LPCS-PZC precursors via this  new method. The advantages of  this process  include:  (1) no  extra solvent  is used;  (2)  the viscosity and preparation time  of LPCS/PZC precursors are adjustable; composition in ZrC-SiC powders  (3)  the  elemental  is  readily  controllable by  varying the LPCS/PZC mass  ratio;  (4)  the  ceramic yield is  relatively high;  (5)  the as-synthesized ZrC-SiC powders exhi bit  spherical morphology with  characteristic  dimensions  of  less  than  60 nm. Finally,  this  hybrid  liquid  precursor  can  also be used for preparing matrix of fiber-reinforced UHTC  composites.  (3)  Oxidation Behavior of ZrC-SiC Powders  To further demonstrate the advantage of the method and the properties of ZrC-SiC composite  superior  powders, oxida tion behavior of as-synthesized powders  is  investigated. The  results are described and discussed in the following sections.  (A)  Thermogravimetric Analyses:  To study the oxida tion  of  the  synthesized  ZrC-SiC  powders  from various  LPCS/PZC mass  ratios,  thermogravimetric  analyses  under  air flow were  carried out  (Fig. 5). For all  samples,  the TG  curves [Fig. 5(a)] are divided into two sections. Below 700°C, at ~400°C, and a maximum weight gain is obtained at ~570°C, and then a weight loss occurs at ~600°C [Fig. 5(a)]. At the same time, the maximum peaks in this stage can also  it  is obvious  to see  that oxidation initiates  be clearly identified in DTG curves  [Fig. 5(b)], which corre spond to the observed rapid weight gain. The first  increase in  mass  and  then  the  decrease after a maximum can be at ~400°C, both ZrC and SiC would  explained as  follows:  show a measured mass gain upon oxidation, and the formed  oxide scales would not volatilize,  so the mass  increases. The  highest percentage of mass gain achieves 5.34%, 5.12%, and  (a)  (b)  (c)  (d)  Fig. 3.  Powder XRD patterns  of  samples  derived  from hybrid  precursors with various LPCS/PZC mass ratios 1550°C in argon: (a) 1:1 (b) 2:1 (c) 3:1 (d) 4:1.  after  pyrolysis  at  Table II.  Elemental Composition and Residual Carbon of ZrC-SiC Powders Derived from Hybrid Liquid LPCS-PZC Precursors  with Various LPCS/PZC Mass Ratios  LPCS/PZC mass ratio  Elemental content  (wt%)  Chemical composition  Residual carbon (wt%)  Zr  Si  C  O  1.0  37.5  38.6  20.47  3.43  ZrSi3.345C4.139O0.520 ZrSi6.445C8.064O0.484 ZrSi9.732C12.220O0.433 ZrSi12.933C16.727O0.359  —  2.0  24.2  48.0  25.74  2.06  1.98  3.0  17.6  52.7  28.36  1.34  3.45  4.0  13.8  54.9  30.43  0.87  5.08  200  Journal of  the American Ceramic Society—Wang et al.  Vol. 98, No. 1  \\x0c', '4.29% when the mass ratio of LPCS/PZC is ~600°C,  1.0,  2.0,  and  4.0,  respectively. At  the  oxidation  of  free  carbon  results in a mass loss. Moreover, the weight gain temperatures are about 568°C, 576°C, and 581°C for the mass ratio of LPCS/PZC = 1.0, 2.0, and 4.0, respectively. The increased  weight  gain temperature ZrC-SiC  indicates  that  the oxidation resis tance  of  powders  gradually  improves  with  the  increased SiC content. 800°C-1400°C region,  In  the  continuous weight  gain  is  observed [Fig. 5(a)] before the completion of oxidation. The  DTG curves [Fig. 5(b)] are also in good agreement with that  of TG curves. In this region, the weight gain ~850°C and terminates at ~1350°C, suggesting that  begins  at  the start ing and ending oxidation 850°C and 1350°C,  temperatures  in  this  section  are  respectively. The final weight gains are the ZrC-SiC  30.51%, 23.62%, and 21.99%,  respectively,  for  powders derived from precursors with mass  ratio of LPCS/  PZC is 1.0, 2.0, and 4.0. The decreased ultimate weight gains  indicate that oxidation resistance significantly improves with  the  increased SiC content. Based on the  chemical  composi tions provided (Table II), the calculated theoretical gains for LPCS/PZC = 1.0 and 4.0 are 33.4% and 36.3%, respectively. that LPCS/PZC = 1.0 sample is ~94% oxidized, This suggest while the LPCS/PZC = 4.0 sample is only oxidized ~61% of  complete  oxidation.  The  theoretical  gains  analyses  also  strengthen the argument  that higher SiC content  is favorable  for better oxidation resistance. All  the above results demon strate that  the oxidation resistance at both intermediate and  high-temperature regions is significantly improved with the increased SiC in the ZrC-SiC composite powders.  (B)  Isothermal Oxidation at  400°C-1300°C:  Figure 6  displays the isothermal oxidation kinetics between 400°C and 1300°C in flowing air,  at  temperatures  from which one  can see that  the isotherms exhibit  two steps,  that  is, a rapid  oxidation step and a slow oxidation step. The first step corre sponds to very rapid oxidation kinetics,  the mass gain in the  first hour varied from 1% to 17% depending on the oxida tion temperature. The second step is a slower oxidation pro cess  with  a  weight  increase  of  less  than  14% for  2-h  treatment in most cases. This two-step manner at temperatures higher than 1000°C. As tabulated in Table II, the estimated theoretical mass gain is ~33.4% to almost fully oxidize into ZrO2 + SiO2 for powders derived from LPCS/PZC = 1.0. In this case, ~12.7% the mass gain is attributed to ZrO2 formation and ~20.7% ascribed to SiO2 formation. These fractions are agreed quite well for the initial step of oxidation in Fig. 6. For high-temperature oxidation above 1000°C,  is  enhanced  of  is  the  it  shows  that  the  entire ZrC fraction oxidizes very rapidly, and followed by a  (a)  (b)  (c)  Microstructure of ZrC-SiC powders Fig. 4. precursor with LPCS/PZC = 1.0 in mass  derived  from hybrid  ratio: SEM micrograph of  typical morphologies  (a), TEM images of an overview observed at  low magnification (b), and a magnified view showing the ZrC and  SiC nanocrystallite (c).  (a)  (b)  Fig. 5. Weight gain (a) and differential thermogravimetry (b) curves of ZrC-SiC powders derived from hybrid precursors with various  LPCS/PZC mass ratios.  January 2015  Synthesis of ZrC-SiC Powders  201  \\x0c', 'slower oxidation of SiC.26,27 Moreover,  it  should be noted  that the ultimate weight gain owing to the oxidation at 550°C is larger than that at 600°C [Fig. 5(a)], suggesting that an anomalous oxidation with high kinetics occurs near 550°C. This anomalous oxidation can be explained as follows. Previous work by Wang et al.28 has demonstrated that  if  the oxidation is mainly controlled by diffusion,  the extent  of oxidation should increase with temperature. However, the increase reaches a maximum around 550°C and then weight declines at ~600°C [Fig. 5(a)], oxidation decreases with temperature. So, in this study, of ZrC-SiC at 550°C obeys  suggesting that  the  extent of  this  anomalous  oxidation  another  oxidation law with high kinetics  rather  than the diffusion controlled law.  It  is also intriguing to find that  the slope of  the isotherm in the first oxidation step for the dized at 1100°C and 1300°C is less than that 1000°C and 1200°C,  samples oxi for the samples  oxidized  at  respectively,  as  shown  in  Fig. 6. The reason will be discussed in detail below. In addi tion,  because  oxidation  is  a  surface  phenomenon, we  also  consider  that  these discrepancies may be within the  experi mental  error  limits of powder oxidation. To further underZrC-SiC  stand  the  oxidation  behavior  of  powders  at  different  temperatures,  the samples after a 2-h oxidation are  characterized by XRD and SEM.  (C)  Phase Evolution  and Morphology Change During  Oxidation:  Figure 7 shows  the XRD patterns of  the pow ders derived from different LPCS/PZC mass ratios of 1.0, 2.0, and 4.0 after oxidation at 600°C [Fig. 7(a)] and 1300°C [Fig. 7(b)], respectively. At 600°C, SiC, ZrO2 (m-ZrO2 and tZrO2), and a few weak peaks of ZrC are clearly identified, indicating that ZrC mainly suffers oxidation at 600°C and  consequently  there  remains  less ZrC in  the  final  oxidized  specimens. The  decrease  of m-ZrO2 with LPCS/PZC is due to the increased C in LPCS, which is usein stabilization t-ZrO2.24,29 Moreover, enter the ZrO2 lattice and substitute O so that of the presence of reported ZrO2-xCy is not excluded here.30,31 Meanwhile, SiC would also oxidize in this temperature but at very slow rates.26 It should also be noted that  the  increase  of  ful  the C atom may  the possibility  the  SiO2 hence  remains  amorphous at this lower temperature and is to XRD. At 1300°C, besides m-ZrO2 and tZrO2, SiO2 is detected, suggesting that both ZrC and SiC subject to further oxidation at 1300°C. Since only SiO2 and ZrO2 phases are formed at 1300°C, the decrease in the slope of the isotherm with temperature should be attributed to the the ZrC-SiC pow invisible  formation rate of oxides on the surface of ders.28 In this case,  the most possible reason for this decrease  with  temperature  is  associated with  the  formation  rate  of  SiO2 instead of ZrO2, because SiO2 than ZrO2. With increasing temperature, protective SiO2 formed more rapidly and makes the oxidation rate slower  is much more protective  is  in  the first steps (see Fig. 6).  From the above results,  the oxidation of ZrC-SiC powders (1) below 700°C, ZrC oxi experiences a two-stage process:  dizes quickly to generate a ZrO2 at very slow rates and thereby the SiO2 remains amorphous; (2) between 800°C and 1400°C, both ZrC and SiC subject  scale, whereas SiC oxidizes  to  oxidation,  the oxidized product consists of ZrO2 and SiO2. Accordingly, the oxidation path of ZrC-SiC powders can be described with the following equations (2)-(4):  2ZrC (s) þ 3O2 (g) ¼ 2ZrO2 (s) þ 2CO (g)  (2)  ZrC (s) þ O2 (g) ¼ ZrO2 (s) þ C (s)  (3)  2SiC (s) þ 3O2 (g) ¼ 2SiO2 (s) þ 2CO (g)  (4)  The morphologies of the ZrC-SiC powders derived from the hybrid precursor with LPCS/PZC = 1.0 after exposure to  air  at  various  temperatures  are  also examined by SEM. the ZrC-SiC particles  It  can be  seen from Fig. 8 that all  are  porous after oxidation and remain intact without  formation  of residue defects 600°C [Fig. 8(a)],  and  cracks. For  the  sample  exposed  at  it  is clearly observed that uniform nanosize  grains distribute on the particle. According to XRD analysis  [Fig. 7(a)],  these small grains are mainly ZrO2 (m-ZrO2 and t-ZrO2) because the SiO2 remains amorphous at this lower temperature. For the sample exposed at 1300°C [Fig. 8(d)],  it  can  be  seen  that  the  sample  shows  grain  growth  and  the  Isothermal weight changes of ZrC-SiC powders Fig. 6. from hybrid precursor with LPCS/PZC = 1.0 in mass to air at 400°C-1300°C as a function of at 550°C is plotted as dotted line.  derived  ratio exposed  time. Anomalous  isotherm  (a)  (b)  Fig. 7.  Powder XRD patterns of ZrC-SiC powders derived from  hybrid precursors with various LPCS/PZC mass 600°C (a) and 1300°C (b) in air for 2 h.  ratios oxidized at  202  Journal of  the American Ceramic Society—Wang et al.  Vol. 98, No. 1  \\x0c', 'oxidation produced SiO2 [Fig. 7(b)] bonds the together and forms agglomerates.27 Furthermore,  particles  the amount  of  voids  obviously  decreases  as  the  oxidation  temperature  increases.  It  is  possible  that  the  reduction  in  voids  is  the  result of  liquid phase sintering due to the molten SiO2. Even the whole powder sample does not sinter, a sintering of the  if  outermost  layer would reduce apparent voids, and the oxi dized scales act as a thick coating that could inhibit oxygen  permeation in the powder sample.  (D)  Oxidation Mechanism:  Below 700°C, the oxidation of ZrC-SiC is dominated by ZrC, because SiC oxidizes at quite slow rates.10,32 Zr is oxidized more easily than carbon.33 The oxidation of ZrC results in the formation of ZrO2 and carbon. Most of the oxidation produced carbon would  be oxidized to generate gas CO. However,  the formation of  ZrO2 tribute to a significant decrease of weight gain rate around 600°C. As an apparent hump between 400°C and a result, 600°C can be observed [Fig. 5(a)], which is in consonance with the previous results of Yan et al.6 and He et al.34 At from 800°C to 1400°C,  scale  can act as an oxygen diffusion barrier and con temperatures  the oxidation of SiC is  predominant.  The formation ZrC-SiC  of  SiO2 with improved  is  protective  and  endows  the  powders  oxidation  resistance.  IV.  Conclusions  In this study, a simple process for synthesizing ZrC-SiC com posite  powders  from hybrid  liquid  precursors  is  reported.  PZC dissolves  directly  into  LPCS without  extra  solvent,  forming hybrid LPCS-PZC precursors with excellent miscibil ity. The viscosity and preparation time of LPCS-PZC precur sors are adjustable, and the elemental composition of  the as synthesized  powders  is  controllable by varying the LPCS/ the hybrid precursors at 1550°C PZC mass ratio. Pyrolysis of results in the formation of ZrC-SiC powders with grain sizes 60 nm. The ZrC-SiC powders  less  than  exhibit  spherical  morphology and the ZrC and SiC phases are uniformly dis tributed.  Investigation  on  the  oxidation  behavior  demon strates that ZrC-SiC  the initial and final oxidation temperatures of from LPCS/PZC = 1.0 powders derived 400°C and 1350°C, precursor are respectively, suggesting good oxidation resistance. Below 700°C, nonprotective ZrO2 scale forms; above 1000°C the oxidized product consists of  the  hybrid  ZrO2 and SiO2, which enables ZrC-SiC with improved oxidation resistance. The oxidized product retains intact without  formation of residue cracks.  Acknowledgments  This work was supported by the National Outstanding Young Scientist Foun dation for Y. C. Zhou under grant no. 59925208, and the Natural Sciences  Foundation of China under grant nos. 50672102, 50832008.  References  1R. Savino, M. De. S. Fumo, D. Paterna, and M. Serpico, “Aerothermody namic Study of UHTC-Based Thermal Protection Systems,” Aerosp. Sci. Technol., 9, 151-60 (2005). 2F. Monteverde, A. Bellosi, and L. 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  "_id": 255,
  "PDF": "Synthesis, densification and characterization of TaB2–SiC composites.pdf",
  "Text": "['Synthesis, densification and characterization of TaB2-SiC composites  Roberta Licheri, Roberto Orru` *, Clara Musa, Giacomo Cao **  Dipartimento di Ingegneria Chimica e Materiali, Unita` di Ricerca del Consorzio Interuniversitario Nazionale per la Scienza e Tecnologia dei Materiali (INSTM),  and Unita` di Ricerca del Consiglio Nazionale delle Ricerche (CNR) Dipartimento di Energia e Trasporti, Universita` degli Studi di Cagliari, Piazza D’Armi,  09123 Cagliari, Italy  Received 4 September 2009; received in revised form 7 October 2009; accepted 21 October 2009  Available online 22 November 2009  Abstract  The TaB2-27.9 vol% SiC composite was synthesized by self-propagating high-temperature synthesis starting from mechanically activated Ta, B4C and Si reactants. The obtained powders were spark plasma sintered at 1800 8C and 20 MPa for 30 min total time, thus obtaining a 96% dense  product. The latter one was characterized in terms of microstructure, hardness, fracture toughness, and oxidation resistance. The obtained results,  particularly the fracture toughness, are promising when compared to those related to analogous materials reported in the literature and fabricated  with similar and different processing routes. # 2009 Elsevier Ltd and Techna Group S.r.l. All  rights reserved.  Keywords: Tantalum diboride; Ball milling; Self-propagating high-temperature synthesis; Spark plasma sintering  1.  Introduction  It  is well known that  the combination of their high melting  temperature, hardness,  electrical  and thermal  conductivities,  and good chemical resistance makes transition metal diborides  such as ZrB2, HfB2, TiB2, TaB2, etc., potential candidates for  high temperature structural applications [1-3]. Among them,  all belonging to the general class of Ultra-High Temperature  Ceramics  (UHTCs),  only  few  investigations  have  been  dedicated  so  far  to  tantalum diboride  based materials,  as  compared to the other UHTC systems. The  effect of TaB2  addition on the characteristics, mainly the oxidation resistance,  of some refractory composites was studied by different authors  [4-7]. Talmy et al. [4] found that oxidation resistance of ZrB2-  SiC materials was  significantly  improved when  partially  replacing ZrB2 with TaB2. More specifically, when comparing  the  effect  of  the  addition  of  10 mol. % of  various metal  diborides, i.e. TiB2, NbB2, VB2, CrB2, and TaB2, the latter one  was found to guarantee the higher protection to the modified  ZrB2-SiC when  exposed  to  air  at  1300 8C for  5 h. More  recently,  the study of the influence of the TaB2 content on the  oxidation resistance of ZrB2-based ceramics received renewed  attention [5,7]. The improvement in the oxidation resistance of  the ceramic composite as a consequence of the addition of TaB2 beyond 3.32 mol.% was confirmed at 1200 and 1400 8C, while 1500 8C. Moreover,  this  property  became worsen  at  the  beneficial  influence  of  tantalum diboride  on  the  oxidation  resistance, densification behaviour and mechanical properties  of TaC-10 wt.% TaB2, as compared to monolithic TaC, was  also demonstrated [6].  In this  study,  tantalum diboride was  synthesized by reducing Ta2O5 with B4C and graphite in a tube  furnace  under Ar  atmosphere,  as  described  in  a  previous  investigation [8], where the resulting powders were subsequently consolidated by HP at 2100 8C to obtain a 98% dense  product.  The preparation of TaB2-TaC composites by combustion  synthesis  starting  from Ta, B4C and  graphite was  recently  studied [9]. Due to the non-sufficiently high enthalpy of the synthesis reaction, the reactants were firstly preheated at 200 8C  to make the process self-sustaining after  ignition.  In  the  present work,  the  fabrication  of  dense  TaB2-  27.9 vol% SiC composite  is  attempted  for  the  first  time  following  a  processing  route  successfully  used  for  the  obtainment of bulk ZrB2-25 vol% SiC [10], ZrB2-40 vol%  www.elsevier.com/locate/ceramint  Available online at www.sciencedirect.com  Ceramics International 36 (2010) 937-941  * Corresponding author. Tel.: +39 070 6755076; fax: +39 070 6755057.  ** Corresponding author. Tel.: +39 070 6755058; fax: +39 070 6755057.  E-mail addresses: orru@visnu.dicm.unica.it (R. Orru` ),  cao@visnu.dicm.unica.it  (G. Cao).  0272-8842/$36.00 # 2009 Elsevier Ltd and Techna Group S.r.l. All rights reserved.  doi:10.1016/j.ceramint.2009.10.028  \\x0c', 'ZrC-12 vol% SiC [11], HfB2-26.5 vol% SiC,  and HfB2-  40.6 vol% HfC-11.2 vol% SiC [12]. This method  basically  consists of firstly obtaining in situ the UHTC composite by self propagating  high-temperature  synthesis  (SHS)  and,  subse quently,  consolidating  it  by  spark  plasma  sintering  (SPS),  where the powders and the die containing them are directly  crossed by a pulsed current  [13].  It  is well established that  heating  processes  are  strongly  accelerated  during  SPS  in  contrast with the relatively long processing times (typically on  the order of hours) encountered in conventional HP.  To promote the SHS character in the synthesis process,  the  mechanical activation by ball milling of the starting reactants is  also considered in this investigation.  2. Experimental materials and methods  The raw materials utilized in the present study were Ta (Alfa-Aesar, \\x00325 mesh, 99.9% purity), B4C (Alfa-Aesar, 1- 7 mm particle size, >99.4% purity) and Si (Aldrich Chemical, \\x00325 mesh, >99% purity) powders. The starting mixture to be  reacted by SHS was prepared by mixing reactants according to  the following stoichiometry:  2Ta þ B4C þ Si ! 2TaB2 þ SiC  (1)  where  the TaB2-27.9 vol% SiC composite, which will  be  indicated as TS in what  follows,  is expected to be formed.  Mechanochemical activation (MA) of reactants was carried  out by co-milling them in a SPEX 8000 (SPEX CertiPrep,  USA)  shaker mill  apparatus with  two  steel  balls  (13 mm  diameter, 8 g weight) for 20 min milling time interval and ball  to powders or charge ratio (CR) equal  to 1.  About 10 g of the MA powders were uniaxially pressed to  form cylindrical pellets with a diameter of 10 mm, height of 30 mm and a green density of \\x1850% of the theoretical value.  The synthesis process was conducted inside a reaction chamber  under Ar atmosphere. The reaction front was generated at one  pellet  end  by  using  a  tungsten  coil maintained  electrically  heated  for  few seconds  until  the  synthesis  reaction was  initiated. The temperature during SHS evolution as well as the  average velocity of the combustion wave was determined using  thermocouples (W-Re, 127 mm diameter, Omega Engineering  Inc.) embedded in the pellet. About 4 g of  the obtained SHS  product  to be densified by SPS was first ground for 20 min by  the mill apparatus described above using a stainless steel vial  with two steel balls (13 mm diameter, 8 g weight). Particle size  distribution of the resulting powders was determined by means  of a laser  light scattering analyser  (CILAS 1180, France).  The densification step of MA-SHS powders was carried out  using an SPS 515 equipment (Sumitomo Coal Mining Co., Ltd.,  Japan).  The  powders  were  first  cold  compacted  into  a  cylindrical  graphite  die  (outside  diameter,  35 mm;  inside  diameter,  15 mm;  height,  40 mm)  lined with  graphite  foils  (0.13 mm thick, Alfa Aesar, Karlsruhe, Germany) to protect the  die  and  facilitate  sample  release  after  synthesis. The  die  containing the sample was then placed inside the SPS reaction  chamber  that was evacuated down to 10 Pa.  The  effect  of  the  total  sintering  time,  tT,  on  sample  densification was investigated by performing all SPS experi ments at constant TD = 1800 8C,  values  of  the  dwell  temperature,  the mechanical pressure, P = 20 MPa, and the  heating time, tH = 10 min, i.e. the time required to reach the TD  value when starting from ambient temperature. For the sake of  reproducibility, each experiment was  repeated at  least  twice.  Further details on the experimental procedure and set-up used  in this work for SHS and SPS can be found in previous works  [10-11,14].  The relative density of the SPS products was measured using  the  Archimedes method.  The  theoretical density 9.98 g/cm3, was  of  the  corresponding  composite,  i.e.  calculated  through rule of mixture [15], by considering the density values of TaB2 and SiC as 12.6 and 3.2 g/cm3,  respectively. Phase  identification was performed by a Philips PW 1830 X-rays  diffractometer using a Ni filtered Cu Ka radiation (l = 1.5405 A˚ ). The microstructure  and local phase composition of end  products were  examined  by  scanning  electron microscopy  (SEM) (mod. S4000, Hitachi, Japan) and energy dispersive X rays  spectroscopy  (EDS)  (Kevex  Sigma  32  Probe, Noran  Instruments, USA),  respectively.  Vickers hardness and fracture toughness (KIC) evaluation of  the  SPSed  products was  performed  using  a  Zwick  3212  Hardness tester machine (Zwick & Co. GmbH, Germany) for  both 3 and 10 kg applied loads with a dwell  time of 18 s.  Oxidation resistance was determined by thermogravimetric  analysis  (TGA) using a NETZSCH (Germany) STA 409PC  Simultaneous DTA-TGA Instrument under 0.1 L/min air flow. Non-isothermal tests, consisting of heating slowly (2 8C/min) the specimen from room temperature to 1450 8C, as well as isothermal runs at 1450 8C for about 4 h, have been performed.  3. Results and discussion  All the attempts carried out to make the synthesis reaction (1)  self-sustaining  failed when  starting  from simply  blended  reactants, in contrast with the classical SHS behaviour displayed  by the analogous binary 2ZrB2-SiC (ZS) and 2HfB2-SiC (HS)  systems,  recently  fabricated  following  the  same  approach  [10,12]. This  outcome  is consistent with the corresponding i.e. \\x00DH o r ¼ 348:364 kJðTSÞ, 647.266 (ZS) and 674.042 kJ (HS) [16], significantly lower for the Ta enthalpies of  reaction,  based composite. Thus,  in the latter case,  the MA of starting  mixture under  the ball milling (BM)  conditions  reported in  Section  2  was  required  to  promote  the  self-propagating  behaviour  in  reaction  (1).  Correspondingly,  the measured  combustion temperature and front velocity were 1850 \\x06 50 8C and 4.5 \\x06 0.5 mm/s, respectively.  equal  to  The need of reactants activation manifested in this work is in  agreement with the behaviour  recently observed during the  preparation of TaB2-TaC by SHS starting from Ta, B4C and  graphite, where the support of the combustion synthesis process was achieved by preheating the reacting pellet at 200 8C [9].  The diffraction pattern of the SHS product obtained in our  study is reported in Fig. 1 along with those of the corresponding  original  reactants before and after  the mechanical activation.  R. Licheri et al. / Ceramics International 36 (2010) 937-941  938  \\x0c', 'No  remarkable  effects  induced  by  the BM treatment  are  evidenced from the XRD results, other  than a  slight peaks  broadening  as  an  indication  of  crystal  size  refinement  and  internal  strain  increase  in  the  processing  powders.  The  enhancement of chemical  reactivity of  starting powders as a  consequence  of  their mechanical  treatment  can  be mainly  ascribed to the interfaces  formation among reactants, which  allows to overcome diffusion limitations.  All major  peaks  corresponding  to TaB2,  as well  as  the  {1 1 1} and {2 2 0} reflections of SiC were detected by XRD.  In addition, no other secondary phases were found in the final  product.  In conclusion,  it  is possible  to state  that  the MA  treatment allows  for SHS to proceed to completion with the  formation of  the desired composite constituents according to  reaction (1).  Once converted in powder  form, SHS products have been  characterized  in  terms  of  particle  size  distribution  and  microstructure. The obtained results  are  shown in Fig. 2(a  and b). Rather fine particle powders are obtained, being about  50% less than 1 mm in size. In particular, all powders have particle size less than 30 mm and d50 = 1.19 \\x06 0.09 mm. The densification of the TS powders was investigated by  examining the effect of the total sintering time (tT) during SPS TD = 1800 8C,  in  the  range  0-30 min,  when  setting  P = 20 MPa,  and  th = 10 min.  The  conditions  above were  chosen on the basis of  the results  found in previous  studies,  where the consolidation by SPS of SHSed 2ZrB2-SiC [10] and  2HfB2-SiC [12] powders was performed.  Typical  sample  shrinkage  and  temperature  time  profiles  recorded during the SPS process are reported in Fig. 3(a and b)  for  the  case  of  tT = 30 min  along with  the  corresponding  electrical behaviour  showing the  current  and voltage mean  values. The most  significant  sample displacement change is  observed to occur after about 6 min from the beginning of the current application, when the temperature is above 1000 8C.  This increase continued until the TD value was achieved, while  no  relevant  changes were  evidenced  during  the  isothermal  stage. As  far  as  the  electrical  behaviour  of  the  system is  concerned (cf. Fig. 3b),  it may be seen that  the current and  voltage are augmented during the non-isothermal heating (0-10  min), to satisfy the selected thermal program. It should be noted  that  the current almost  reached the maximum value (1500 A)  allowable with the SPS apparatus used in the present work.  Afterwards,  first  rapidly,  then  smoothly,  both  parameters  decrease down to the corresponding stationary mean values, i.e.  about 1300 A and 5.5 V,  respectively.  The situation described above also includes the experiments  carried out when the SPS process duration was  shorter  than  30 min, while keeping the  same  all  the other  experimental  parameters. Thus,  the considerations made here hold also true  in these cases. Regarding the corresponding SPSed products  density, as shown in Fig. 4, it increases from about 91% of the  theoretical value at the end of the non-isothermal heating stage up to about 96% when maintaining the sample at 1800 8C for 20  more minutes.  Two back-scattered SEM micrographs at different magni fications of the 96% dense product are reported in Fig. 5(a and  b). A rather homogeneous and fine microstructure consisting of  TaB2 (brighter) and SiC (darker) phases well distributed all over  the sample is obtained.  The result obtained by TGA in the temperature range 700- 1450 8C during the non-isothermal oxidation test  is compared  in Fig. 6 with that of monolithic TaB2 reported in literature [6].  While mass gain monotonically increases with temperature in  the  latter  case,  a  relative maximum is  displayed  at  about  Fig. 1. Comparison of XRD patterns of  (a) original  reactants,  (b) mechan ochemically activated reactants and (c) products obtained by MA-SHS accord ing to reaction (1).  Fig. 2. Size distribution (a) and SEM back-scattered micrograph (b) of TaB2-  SiC SHS powders after 20 min ball milling.  R. Licheri et al. / Ceramics International 36 (2010) 937-941  939  \\x0c', '940  R. Licheri et al. / Ceramics International 36 (2010) 937-941  Fig. 3. SPS outputs temporal profiles recorded during the fabrication of dense  borosilicate protective layer, as discussed in details by several  Fig. 5. SEM back-scattered micrographs of dense 2TaB2-SiC product: 500\\x02 and (b) 5000\\x02.  (a)  2TaB2-SiC: (a) temperature and sample shrinkage, (b) mean current and mean voltage (TD = 1800 8C,  tT = 30 min, P = 20 MPa).  tH = 10 min,  intensity  1200 8C by the two-phases material. Such oxidative behaviour  studies reported in the literature on this subject  [5,17-18].  In  particular, it is apparent the role played by SiC in protecting the  TS material from oxidation at temperatures in the range 1350- 1450 8C, since the weight gain of pure TaB2 becomes more than  is  also typically exhibited by other MB2-SiC (M = Zr, Hf)  twice with respect to the TS system and tends to increase in an  systems  reported in the literature,  including the ZS and HS  exponential manner. The fact that the composite material shows  products  [10,12]  obtained  using  the  same  processing  route  a  normal weight  gain  utilized in this work, except for the mechanical activation. This  temperatures  lower  than  slightly higher 1350 8C may  than  pure  TaB2  at  be  ascribed  to  the  feature is consistent with the presence of SiC in the composite,  different relative densities of the two samples, i.e. about 96 and  being  the  latter  one  responsible  for  the  formation  of  a  98% for  the binary and single phase systems,  respectively.  Fig. 4. Effect of SPS time on relative density of sintered 2TaB2-SiC powders synthesized by MA-SHS (TD = 1800 8C,  tH = 10 min, P = 20 MPa).  Fig. 6. Comparison among specific weight changes as a function of temperature  during non-isothermal TGA oxidation in air of TaB2 and 2TaB2-SiC dense  products.  \\x0c', 'The measured Vickers  hardness  and  fracture  toughness  properties of  the best TS product obtained in this work are  reported in Table 1 along with the corresponding values related  to the ZS, HS, ZZS and HHS materials fabricated following the  same  route.  For  the  sake  of  comparison,  the mechanical  properties reported in the literature for dense TaB2 and TaC-  10 wt.% TaB2  are  also  included  in  Table  1. As  already  evidenced  in  previous  papers  [10-12],  properties  of  ultra refractory  products  obtained  by  combining  SHS  and  SPS  techniques are generally among the best, when compared to  those reported in the literature for similar materials produced by  other, generally more time and energy-consuming, competitive  methods. The combination of  the in situ synthesis  (SHS) of  composite powders with a rapid densification technique (SPS)  was considered as responsible of the observed advantages. This  conclusion is  still valid in this  study. Furthermore,  it  is seen  from Table 1 that  the fracture toughness value (KIC) of the TS  ceramic is even higher  than ZS and HS.  Such improvement  is  likely associated to the mechanical  treatment provided in this work to the reactants before the SHS  process  that,  other  than  increasing  powder  reactivity  to  guarantee the SHS character of the synthesis reaction, produced  powder refinement as well as a better phases distribution in the  composite.  This  aspect  leads  to  a  relatively  finer  and  homogeneous  microstructure  in  the  final  material  and,  consequently,  to an improvement of its mechanical properties.  4. Conclusions  A near  fully  dense TaB2-27.9 vol% SiC composite was  fabricated by SPS starting from powders synthesized by SHS.  The mechanical activation for 20 min (CR = 1) of the Ta, B4C  and Si reactants was required to make the reactive process self sustained and obtaining in situ the desired ceramic constituents.  In addition, when the resulting powder products were SPSed for total processing time of 30 min by setting TD = 1800 8C,  a  P = 20 MPa,  tH = 10 min, a 96% dense material  is obtained.  The  characterization  of  the  composite  revealed  that  it  possesses fracture toughness values better  than the analogous  UHTCs reported in the literature, included the ZrB2and HfB2 products fabricated following the same route, with the exception  of the mechanical treatment. Regarding the oxidation behaviour, a significant improvement at temperature higher than 1350 8C is  observed as compared to the monolithic TaB2, as a consequence  of the beneficial presence of SiC in the composite material.  Acknowledgments  IM (Innovative Materials) S.r.l., Italy, is gratefully acknowl edged for granting the use of SPS apparatus. The authors thank  Eng. Leonardo Esposito (Centro Ceramico di Bologna,  Italy)  for performing hardness and fracture toughness measurements.  References  [1] K. Upadhya,  J.M. Yang, W.P. Hoffmann, Materials  for ultrahigh tem perature structural applications, American Ceramic Society Bulletin 58  (1997) 51-56.  [2] W.G. Fahrenholtz, G.E. Hilmas,  I.G. Talmy,  J.A. Zaykoski, Refractory  diborides  of  zirconium and  hafnium,  Journal  of American Ceramic  Society 90 (2007) 1347-1364.  [3] R. Rapp, Materials for extreme environments, Materials Today 9 (5) (2006) 6.  [4]  I.G. Talmy, J.A. Zaykoski, M.M. Opeka, S. Dallek, Oxidation of ZrB2  ceramics modified with SiC and group IV-VI transition metal diborides,  Electrochemical Society Proceedings 12 (2001) 144-158.  [5] F. Peng, R.F. Speyer, Oxidation resistance of fully dense ZrB2 with SiC,  TaB2, and TaSi2 additives, Journal of American Ceramic Society 91 (5)  (2008) 1489-1494.  [6] X. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Densification, mechanical  properties, and oxidation resistance of TaC-TaB2 ceramics, Journal of the  American Ceramic Society 91 (12) (2008) 4129-4132.  [7] F. Peng, Y. Berta, R.F. Speyer, Effect of SiC, TaB2 and TaSi2 additives on  the isothermal oxidation resistance of  fully dense zirconium diboride,  Journal of Materials Research 24 (5) (2009) 1855-1867.  [8] X. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Synthesis, densification, and  mechanical properties of TaB2, Material Letters 62 (2008) 4251-4253.  [9] C.L. Yeh, Y.L. Chen, An experimental study on self-propagating high temperature  synthesis  in  the Ta-B4C system,  Journal  of Alloys  and  Compounds 478 (1-2) (2009) 163-167.  [10] R. Licheri, R. Orru` , A.M. Locci, G. Cao, Efficient  synthesis/sintering  routes  to obtain fully dense ZrB2-SiC ultra-high-temperature  ceramics  (UHTCs), Industrial Engineering Chemistry Research 46 (2007) 9087-9096.  [11] R. Licheri, R. Orru` , C. Musa, G. Cao, Combination of SHS and SPS  techniques  for  fabrication  of  fully  dense ZrB2-ZrC-SiC composites,  Material Letters 62 (2008) 432-435.  [12] R. Licheri, R. Orru` , C. Musa, A.M. Locci, G. Cao, Consolidation via spark  plasma  sintering of HfB2/SiC and HfB2/HfC/SiC composite powders  obtained  by  self-propagating  high-temperature  synthesis,  Journal  of  Alloys and Compounds 478 (2009) 572-578.  [13] R. Orru` , R. Licheri, A.M. Locci, A. Cincotti, G. Cao, Consolidation/  synthesis  of materials  by  electric  current  activated/assisted  sintering,  Materials Science and Engineering R Reports 63 (4-6) (2009) 127-287.  [14] A. Cincotti, R. Licheri, A.M. Locci, R. Orru` , G. Cao, A review on  combustion synthesis of novel materials: recent experimental and model ing results, Journal of Chemical Technology and Biotechnology 78 (2-3)  (2003) 122-127.  [15] F.L. Matthews, R. Rawlings, Composite Materials: Engineering  and  Science, Chapman and Hall, UK, 1994.  [16]  I. Barin, Thermochemical Data of Pure Substances, Wiley-VHC, New  York, 1989.  [17] W.C. Tripp, H.H. Davis, H.C. Graham, Effect of a SiC addition on the  oxidation of ZrB2, American Ceramic Society Bulletin 52 (8) (1973) 612-  616.  [18] F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC  composite, Journal of the European Ceramic Society 25 (7) (2005) 1025-  1031.  Table 1  Properties of dense ceramic composites.  System  Relative  density [%]  Hardness [GPa]  (applied load)  KIC  [MPa m1/2]  Reference  ZS  99.6  16.7 \\x06 0.4 (1 kg) 20.55 \\x06 0.8 (3 kg) 19.2 \\x06 0.6 (10 kg) 16.9 \\x06 0.2 (10 kg) 17.7 \\x06 1.5 (3 kg) 18.3 \\x06 1.1 (10 kg) 25.6\\x06 0.7(0.5 kg) 19.4 \\x06 0.6 (0.5 kg) 19.4 \\x06 0.9 (3 kg) 18.9 \\x06 0.4 (10 kg)  5.0 \\x06 0.3  [10]  HS  >99.9  [12]  7.0 \\x06 0.7 5.9 \\x06 0.5  ZZS  98.7  [11]  HHS  98.5  [12]  6.2 \\x06 0.7 4.5 \\x06 0.3 3.4 \\x06 0.1 8.2 \\x06 0.6 8.4 \\x06 0.8  TaB2  98  [8]  TaC-TaB2  98.6  [6]  TS  96  [This work]  R. Licheri et al. / Ceramics International 36 (2010) 937-941  941  \\x0c']"
},{
  "_id": 256,
  "PDF": "Synthesis, Sintering, and Oxidative Behavior of HfB2–HfSi2 Ceramics.pdf",
  "Text": "['Article  pubs.acs.org/IECR  Synthesis, Sintering, and Oxidative Behavior of HfB2−HfSi2 Ceramics  Clara Musa, Roberta Licheri, Roberto Orru ̀,* and Giacomo Cao  Ingegneria Meccanica, Chimica e dei Materiali, Centro Studi sulle Reazioni Autopropaganti (CESRA), Unita ̀ di Dipartimento di Ricerca del Consorzio Interuniversitario Nazionale per la Scienza e Tecnologia dei Materiali (INSTM), Universita ̀ degli Studi di Cagliari, via Marengo n. 3, 09123 Cagliari, Italy  (1 − x)HfB2 −xHfSi2 ABSTRACT: The self-propagating high-temperature synthesis (SHS) of (x from 0 to 1) ceramics is successfully performed in this work starting from elemental reactants. The presence of HfSi2 in the composite system progressively reduces the exothermic character of the synthesis reaction. Bulk ceramics are then obtained after consolidation by spark plasma sintering (SPS) of the SHS crashed products. It is found that the HfB2 −xHfSi2 composite powders synthesized in a single step by SHS require milder sintering conditions as compared to the mixtures consisting of the ceramic constituents obtained separately by the same route. In addition, 15 vol % is the minimum percentage of HfSi2 required to achieve the complete densification of the starting powders, under the SPS conditions investigated in the present work (I = 1350 A, P = 50 MPa, 30 min total time). Thermogravimetric analysis experiments performed in air flow up to 1450 °C clearly indicate that the introduction of HfSi2 plays a beneficial role for protecting the obtained material from oxidation.  1.  INTRODUCTION  It is well-known that HfB2 represents, along with ZrB2, the most important base material in the class of ultrahightemperature-ceramics (UHTCs) because of its interesting properties, such as high melting point (3380 °C), high thermal (104 W/mK) and electrical (9.1 × 106 S/m) conductivities, and (6.3 × 10−6 K−1).1 These thermal expansion coefficient and other attractive characteristics makes HfB2-based ceramics ideal candidates in various application fields where harsh environments have to be withstood, as for the fabrication of thermal protection systems of components in the aerospace industry (leading edges, nose tips).1 HfB2 has been also recognized as suitable material for cutting tools, refractory linings, microelectronics, and neutron absorbers.1,2 One of the main concerns related to the fabrication and utilization of this family of materials relates to the low intrinsic sinterability of HfB2 powders so that severe consolidation temperature and applied pressure conditions are needed to obtain nearly full dense products. It is well established that the spark plasma sintering (SPS) technology, where the powders to be consolidated and/or the mold containing them are rapidly crossed by an electric pulsed current, offers a powerful tool to overcome the drawback above, since relatively milder sintering conditions are needed in (HP) methods.3 comparison with classical hot pressing Specifically, several bulk HfB2-based ceramic materials, either monolithic or composite, requiring relatively lower sintering temperatures and few minutes holding times, have been fabricated so far by SPS.3−6 Nevertheless, it should be also noted that monolithic HfB2 is characterized by relatively modest resistance to oxidation at temperatures.6 high This fact clearly limits its possible application under such conditions. Several studies reported in the literature have shown that the introduction of suitable amounts of Si-containing phases, such as SiC,4,6−8 MoSi2,5 TaSi2,5 into HfB2 matrices, enhances the  oxidation resistance of the resulting composite material, other than facilitating powder consolidation. Along these lines, one of the ceramics that might be used in combination with HfB2 to improve its resistance to oxidation is represented by hafnium disilicide (HfSi2). Owing to its interesting properties, the latter one is suitable for protection of refractory metals and alloys, electrodes, and electronic devices.9 Despite its importance, HfSi2 is, among the silicide phases, one of the less investigated systems. Indeed, only a few studies have addressed this material either in its monolithic9−11 or composite11−13 forms. For instance, this ceramic compound was prepared by mechanical alloying after 30 h ball milling of ratio.10 More Hf and Si powders in stoichiometric recently, HfSi2 was synthesized from elemental powders using classical furnace methods under flowing argon.11 The electrochemical synthesis of hafnium disilicide starting from a molten salt of NaCl−KCl−NaF−K2HfF6 −K2SiF6 was also accomplished.9 As far as composite materials are concerned, a plate of HfB2 7 vol % HfSi2 was prepared by HP to be subsequently coupled with HfSi2 and HfO2 layers in order to eva luate the ceramic−ceramic the Hf−Si−O compatibility of phases in system.12 The consolidation of HfB2 −5 vol % HfSi2 powders by HP was also performed by Monteverde11 starting from commercial HfB2 powders and the disilicide phase, which was previously synthesized from elemental reactants as described above. More recently, HfC + 3Si mechanically activated powders were reactively sintered by high-frequency induction −S iC heated combust ion synthes is to obta in the H fS i2 composite.13 In the present work, the fabrication by SPS of dense HfB2 xHfSi2 ceramics is investigated for the first time using ceramic  −  −  Special Issue: Massimo Morbidelli Festschrift  Received:  Revised:  Accepted:  Published:  October 1, 2013  January 7, 2014  January 7, 2014  January 7, 2014  © 2014 American Chemical Society  9101  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  Downloaded via FLORIDA INTL UNIV on January 15, 2020 at 21:45:45 (UTC).See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.\\x0c', 'Industrial & Engineering Chemistry Research  powders prepared by self-propagating high-temperature synthesis (SHS).14 In this regard, it should be noted that the combination of SHS and SPS routes was successfully adopted −SiC and in previous studies for the acquisition of dense MB2 −MC−SiC (M = Zr, Hf, or Ta) products.4,15−17 MB2 In this work, the composite mixtures are obtained by blending HfB2 and HfSi2 powders synthesized separately by SHS from elemental reactants. Alternatively, the SHS route is also exploited to obtain the composite product in a single step. Possible benefits in term of resistance to oxidation deriving from the introduction of HfSi2 on hafnium diboride matrix are evaluated by comparing the behavior of the obtained nearly dense composite materials when exposed to high temperatures in air flow during thermogravimetric analysis. It should be noted that, to the best of our knowledge, no studies have been addressed so far in the literature on such a characterization −HfSi2 system. relatively to the HfB2  2. EXPERIMENTAL MATERIALS AND METHODS  SHS experiments were carried out starting from mixtures prepared according to the following stoichiometries:  Hf  Hf  +  +  2.2B  →  HfB2  2Si  →  HfSi 2  (1  +  α  )Hf  +  2.2B  +  α  2  Si  →  HfB  2  +  α  HfSi  2  (1)  (2)  (3)  Commercially available hafnium (Alfa-Aesar, <44 μm, >99.6% purity), silicon (Aldrich, <44 μm, 99% purity) and amorphous boron (Aldrich, <9 μm, 95−97% purity) were used as starting powders. The use of an excess of boron is required to compensate its partial loss during the occurrence of the synthesis reaction.6 To prepare the desired products, reactions 1 and 2 are performed separately. Alternatively, the two reactions above are carried out simultaneously according to −xHfSi2 reaction 3 to synthesize in situ the composite HfB2 material, where the x value was changed in the range of 0−30 vol % (α = 0−0.26). Powder mixing was carried out for 15 min using a SPEX mixer mill 8000 (SPEX CertiPrep, USA) shaker mill, plastic bottle, and zirconia balls milling media. A certain amount of the resulting mixture, in the range of 7− 9 g, depending upon the system to be synthesized, was first uniaxially pressed to form cylindrical pellets (11 mm diameter, 22 ± 2 mm height, and about 50% green density). The pellets were then reacted by SHS to produce the desired ceramic, −xHfSi2)SHS. The namely (HfB2)SHS, (HfSi2)SHS, and (HfB2 synthesis reaction was performed in a closed vessel under argon environment. The detailed description of SHS experiments including the complete setup can be found elsewhere.18 Specifically, a C-type thermocouple (Omega Engineering Inc., USA) embedded into the green compact was typically utilized for temperature measurement during reaction evolution. On the other hand, when thermal levels exceeded the limit of about 2300 °C for C-thermocouples, a two-color pyrometer (Ircon Mirage OR 15-990, USA) was instead used for this purpose. The SHS process gives rise to the formation of porous products. Thus, to obtain the powders to be sintered, about 4 g of the SHS material were crushed for 20 min using the SPEX mixer mill 8000, stainless steel vial and a steel ball (13 mm diameter, 8 g weight). When (HfB2)SHS and (HfSi2)SHS were synthesized separately according to reactions 1 and 2, respectively, the composite  Article  mixture to be processed by SPS was obtained by combining them in appropriate proportions to produce the desired system −x(HfSi2)SHS. composition, hereafter indicated as (HfB2)SHS A Spark Plasma Sintering apparatus (SPS 515 Sumitomo Coal Mining Co. Ltd., Japan) under vacuum (20 Pa) conditions was utilized to consolidate SHS powders. This machine combines a DC pulsed current generator (10 V, 1500 A, 300 Hz), to provide an electric current through the processing powders (6 g) and/or the graphite cylinders containing them, with an uniaxial press (max 50 kN) for the simultaneous application of a mechanical load through the punches. SPS experiments were generally carried out using the graphite cylinder configuration, hereafter indicated as A and characterized by external diameter (De) and inside diameter equal to 35 mm and 15 mm, respectively. Alternatively, a different die configuration, hereafter indicated as B, which has smaller De value (30 mm) and the same inner diameter, was also adopted. Both the dies and related plungers were composed of AT101 graphite (Atal s.r.l., Italy). A prescribed electric current cycle was applied to the SPS system during sintering experiments. This procedure was chosen, instead of the temperature control mode, for the sake of comparison with recent studies reported in the literature on the fabrication of monolithic HfB2.6 Specifically, the maximum value of the mean electric current (I = 1350 A) was achieved in 10 min (tH) and maintained constant for additional 20 min (tD). The mechanical pressure (P = 50 MPa) was held constant during the entire SPS run. The most relevant SPS parameters, temperature, current, voltage between the machine electrodes, mechanical load, and vertical sample displacement (δ), were recorded in real time. The temporal evolution of the parameter is important, as it provides an indication of compact densification, although the thermal expansion of the electrodes/ spacers/die/plungers/sample ensemble also contributes to the measured value.19 Temperature−time profiles were obtained by means of a two-color pyrometer (Ircon Mirage OR 15-990, USA) focused on the lateral surface of the graphite die. To limit heat losses by thermal radiation, thus reducing thermal gradients in radial direction, a graphite felt layer (Atal s.r.l., Italy) was placed around the die, while a graphite foil liner (Alfa Aesar, 0.13 mm thick, 99.8% purity) was used to facilitate product release after sintering. The final density of ceramics was determined on polished dense samples through geometric/gravimetric measurements and using the Archimedes method. In this regard, the theoretical densities of the monolithic and composite products were calculated by considering the density values of HfB2, and HfSi2 as 11.18 and 7.97 g/cm3, respectively. A Philips PW 1830 X-ray diffractometer using Cu Kα radiation (λ = 1.5405 Å) and Ni filter was used for phase identification. Particle size distribution was determined by means of a laser light scattering analyzer (CILAS 1180, France). The microstructure of sintered materials was examined by scanning electron microscopy (SEM) using a Zeiss EVO LS15 microscope equipped with energy dispersive X-rays spectroscopy (EDS), Oxford X-MAX Probe. The oxidation resistance of bulk ceramics was evaluated by performing thermogravimetric analysis (TGA) measurements in air flow (100 mL/min) under isothermal conditions (1450 °C) using a NETZSCH STA 409PC Simultaneous DTA-TGA instrument.  δ  9102  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', 'Industrial & Engineering Chemistry Research  3. RESULTS AND DISCUSSION  3.1. Powders Synthesis. The reactions 1−3 investigated in −xHfSi2 this work for the synthesis of HfB2, HfSi2, and HfB2 systems exhibited, upon ignition, a self-sustaining behavior. A typical example of the temperature profile measured during the evolution of the SHS process is shown in Figure 1 for the  Article  Figure 2. Effect of the HfSi2 content on the maximum combustion temperature (TC) and front velocity (vf) during the SHS process evolution, according to reactions 1−3.  Figure 1. Temperature−time profile recorded during the synthesis of −30 vol % HfSi2 by SHS. HfB2  −30 vol % HfSi2 system. case of the HfB2 It is observed that, as soon the combustion front approached to the position where the thermocouple was embedded into the reacting pellet, the temperature level increased sharply from about 100 °C to its maximum value of approximately 2100 °C in less than 1 s. Specifically, the corresponding heating rate is higher than 105 °C/min. Afterward, the sample cooled down still rapidly, albeit with a relatively lower rate (on the order of 103−104 °C/min) as compared to the heating step, to achieve temperature levels below 1000 °C in about 10 s. The dependence of the combustion temperature and front velocity on the HfSi2 content in the composite system is reported in Figure 2. Specifically, the measured average reaction front velocity and maximum combustion temperature for the single phase system HfB2 were about 8 mm/s and 2700 °C, respectively, and approximately 1.5 mm/s and 1600 °C, respectively, when synthesizing HfSi2. From Figure 2 it is also clearly seen that the presence of HfSi2 progressively reduces both TC and vf parameters, thus indicating that the exothermic character of the reaction for the synthesis of the composite materials is correspondingly mitigated. This outcome is clearly expected if the high enthalpy of HfB2 formation, that is, −ΔHr o = 335.975 kJ/mol20 is compared with the significantly lower heat of reaction relative to the disilicide phase, that is, 69.7 kJ/mol.21 The composition of the materials obtained by SHS can be seen in Figure 3, where the XRD patterns relative to the end resulting from reactions 1−3 are shown. Depending products upon the system under consideration, HfB2 and/or HfSi2 are the only phases detected in the end products through this analysis, while no traces of the original reactants or secondary  Figure 3. XRD patterns of SHS products obtained according to (a) reaction 1, (b) reaction 2, (c) reaction 3 with α = 0.26 equivalent to (HfB2 −30 vol % HfSi2)SHS, and (d) (HfB2)SHS −30 vol % (HfSi2)SHS from ceramic products synthesized separately.  products can be found. For the sake of comparison, the XRD pattern relative to the composite mixture obtained after blending HfB2 and HfSi2 synthesized separately is also reported −30 vol % HfSi2 system. No in Figure 3 for the case of the HfB2 significant differences can be found with respect to the product with the same nominal composition prepared in a single step according to reaction 3. Analogous results are obtained when  9103  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', 'Industrial & Engineering Chemistry Research  analyzing the other composite materials investigated in this work (not shown). Therefore, it is possible to conclude that all the SHS reactions went to completion with the formation of the desired phases. In this regard, it is important to emphasize the fact that the complete reactants conversion was achieved by SHS in less than 20 s, even when processing the less exothermic HfSi2 system. This is an important outcome if compared to the more severe processing conditions (1150 °C end temperature, °C/min heating 10 time, and 45 min dwell time) required when the synthesis of this silicide phase starting from elemental furnace methods.11 powders was addressed using classical As described in the previous section, the SHS porous products were finely crushed by ball milling before being processed by SPS.  3.2. Spark Plasma Sintering of SHS Powders. The SPS  conditions (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min) applied for the consolidation of powders prepared by SHS have been chosen in agreement with results recently reported in the literature.6 Specifically, a 89% dense additive free HfB2 material was obtained from SHS powders when using the processing conditions above. Correspondingly, the maximum temperature °C. The measured was about 1800 influence of the HfSi2 addition on the density of the SPS products is shown in Figure 4. In particular, the results relative to powders obtained after  Figure 4. Effect of the HfSi2 content on the density of bulk products obtained by SPS (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min, De = 35 mm) when starting from composite powders synthesized by SHS in one or two steps.  mixing (HfB2)SHS and (HfSi2)SHS products are compared to the case of the composite material synthesized in a single step by SHS.  Article  −  The general comment is that the presence of HfSi2 is beneficial in both circumstances, particularly as the additive 8−10 content is above vol % to improve HfB2 product densification. In addition, a major effect is clearly displayed −xHfSi2)SHS powders. when processing (HfB2 To verify if the observed behavior could be explained by particle size effect, the granulometry of the starting mixtures to be sintered have been analyzed. The obtained results are −5 vol % HfSi2 summarized in Table 1 for the case of the HfB2 composite material. These data indicate that the mixture obtained with (HfB2)SHS and (HfSi2)SHS powders is characterized by relatively finer particles as compared to the (HfB2 xHfSi2)SHS system. Analogous outcomes are obtained when the comparison is extended to the other compositions. Therefore, particle size features cannot be apparently considered to justify the relatively higher sintering ability exhibited by the composite system synthesized in one step by SHS. It should be noted that a similar finding was reported in the literature16 when investigating the consolidation by SPS of −40 vol. %ZrC−12 ZrB2 vol. % SiC products. Specifically, higher densification levels were achieved, under the same SPS conditions, when using composite powders obtained in a single stage by SHS instead of a mixture of relatively finer commercial ZrB2, ZrC, and SiC materials. One of the possible motivation al.16 indicated by Licheri et to justify such behavior was the higher defect concentration present in SHS powders, with respect to analogous products prepared by alternative methods. This feature is a consequence of the extreme heating and cooling rates conditions (up to 200000 K/min) typically established during reaction front propagation.21 This behavior is confirmed by the temperature data plotted in Figure 1, as discussed previously. However, as far as the two composite mixtures compared in the present work are concerned, the defect concentration mentioned above is not expected to produce differences in ceramic powders behavior, since both of them are synthesized by SHS. On the other hand, the improved intimate contact with the formation of strong interfaces among the different ceramic phases formed in situ during the synthesis of composite materials by SHS likely plays a promoting role in this circumstance. On the basis of these considerations, the different sintering −xHfSi2 system can behavior observed for the case of the HfB2 be apparently due to the reduction of diffusion distances among the two ceramic phases simultaneously produced by SHS, which improves the sintering ability of the resulting powders. Another important feature observed in Figure 4 is that the minimum HfSi2 amount required to achieve full product densification is 15 vol %. Indeed, when the HfSi2 present as sintering aid is equal or lower than 10 vol %, the residual porosity remaining in the final product obtained under these SPS experimental conditions (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min) from such powders is still high. It should be noted that, although the maximum allowable current intensity  −5 vol %(HfSi2)SHS and (HfB2 −5 vol % HfSi2)SHS Powders Table 1. Comparison of Particle Size Distribution of (HfB2)SHS Obtained by SHS in Two or One Step, Respectively, after Receiving a 20 min Ball Milling Treatment  −5 vol %(HfSi2)SHS (HfB2)SHS −5 vol % HfSi2)SHS  (HfB2  d10 (μm)  0.41 5.14 ± 0.05  d50 (μm) 3.61 ± 0.14 5.28 ± 0.24  d90 (μm) 14.64 ± 0.33 21.94 ± 3.08  9104  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', '(BSE) images obtained when examining the 95% dense HfB2 a−c, 10 vol % HfSi2 material are reported in Figures 6 panels  −  Article  Industrial & Engineering Chemistry Research  with the SPS apparatus adopted in the present work is 1500 A, the use of electric currents higher than 1350 A to improve powders densification was not convenient for safety reasons (die/plungers breakage, etc.). This aspect should be also taken into account when the applied pressure was increased, at these current/temperature levels, to values higher than 50 MPa. It should be noted that all the SPS experiments described above are performed using the configuration A (De = 35 mm) of the graphite die. Thus, another possible option to improve product density is to use a die with a relatively smaller external diameter (configuration B); that is, De = 30 mm. Indeed, when the electric current passes through a die with relatively smaller cross section, the corresponding current density increases and so does the electrical power supplied by the Joule effect to the powders undergoing sintering. This possibility was explored by processing the (HfB2 −10 vol % HfSi2)SHS system, and the results obtained using the two die configurations mentioned above are shown in Figure 5. As expected, the densification  −10 vol Figure 5. Comparison of end-product densities in the (HfB2 % HfSi2)SHS system sintered by SPS (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min) using two die configurations: (A) De = 35 mm and (B) De = 30 mm.  level achieved is improved, although the complete densification was not reached, the final product being about 95% dense. This result is consistent with the increased temperature levels, that is, °C higher, about 60 correspondingly measured. A further decrease of the De value, while maintaining all the other condition unchanged, was not possible as it would lead to die breakage. In conclusion, on the basis of the results described above, the consolidation of HfB2 −xHfSi2 powders by SPS is more convenient when starting from composite products synthesized by SHS in a single step. In addition, 15 vol % is the minimum percentage of additive required to achieve the complete densification of the starting powders, under the SPS conditions investigated in the present work (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min, De = 35 mm). 3.3. Sintered Ceramics. The obtained bulk products have been investigated by SEM. Specifically, three backscattered  −10 vol % HfSi2 Figure 6. SEM images of the nearly full dense HfB2 sample obtained by SPS (I = 1350 A, P = 50 MPa, tH = 10 min, tD = 20 min, De = 30 mm): SE (a) and BSE (b) general views, detailed views and related EDS analyses (c,d).  along with the corresponding EDS analysis results. First of all, the presence of residual porosity distributed throughout the sample is confirmed in Figure 6a. In addition, EDS analysis indicates that the hafnium silicide phase is present between HfB2 grains (cf. Figure 6b), whose average size is about 10 μm. Furthermore, a silicon-rich phase can be also found in the sample, particularly localized inside the pores (cf. Figure 6c). The results described above can be compared with the literature,11 experimental findings recently reported in the where the full consolidation of HfB2 −5 vol % HfSi2 powders by HP was attained at 1600 °C, 30 MPa, and 15 min holding time. Although powder densification was performed at temperatures lower than the melting point of HfSi2 (about 1700 °C),22 several features were evidenced in Monteverde’s11 work to indicate that the process proceeded through liquid phase sintering mechanism. In this regard, it was assessed that this mechanism was originated by the reaction of HfSi2 with oxide impurities (B2O3, HfO2) present on the surface of HfB2 particles. In particular, the occurrence of reaction  5HfSi  2  +  2B O  2  3  →  2HfB  2  +  10Si  +  3HfO  2  (4)  was postulated. This assumption was supported by the formation of Si based pockets in the dense product while the corresponding HfSi2 content was drastically reduced after sintering. The presence of Hf−silicide as well as the Si-rich  9105  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', 'Industrial & Engineering Chemistry Research  phase is consistent with the experimental outcomes obtained in the present work, as evidenced in Figures 6 by EDS analysis. On the other hand, relatively higher sintering temperature and larger percentage of HfSi2 had to be adopted in our investigation to achieve the complete powder consolidation in comparison with those required in Monteverde’s11 study. In this context, as reported in the Introduction, it should be mentioned that another study was also addressed in the −HfSi2 composite.12 literature for the sintering of HfB2 Although the amount of HfSi2 (7 vol %) was higher with respect to the investigation by Monteverde,11 relatively more drastic conditions (2000 °C, 27.6 MPa, and 1 h holding time) were applied to prepare a plate of the composite material by HP. Thus, the conditions adopted by Bronson et al.12 seem to be more consistent with those ones required in the present investigation (cf. Figure 4). The discrepancies observed in the three studies cited above might be likely explained on the basis of the different powders characteristics used as starting material. Specifically, as far as the present investigation is concerned, it is possible that a lower amount of oxide impurities is present on the surface of HfB2 powders produced by SHS, as compared to the commercial starting material by Monteverde.11 Thus, on the ones used as basis of the considerations made above, it is likely that in the latter study the occurrence of reaction 4 is favored and as is the formation of a liquid phase at lower temperatures to promote powder sintering. A liquid phase sintering mechanism is expected to occur also in our study, since the temperature (about 1800 °C) achieved during SPS is higher than the melting point of the disilicide phase. In addition, the occurrence of reaction 4 may be invoked to justify the presence of the Si-rich phases evidenced in Figure 6c, although no HfO2 was detected in our samples. effect of the presence of HfSi2 on the oxidation behavior of the HfB2-based ceramics was evaluated using TGA by monitoring the mass change of the sample subjected to an oxidizing environment (100 mL/min air flow) up to 1450 °C. In particular, the measurement recently reported in the literature6 during the isothermal test at 1450 °C for the 98.8% dense monolithic material obtained by reactive SPS is compared in Figure 7 with those ones obtained in the −10 vol present work relatively to two composite systems, HfB2 % HfSi2 (95% dense sample) and HfB2 −15 vol % HfSi2 (fully dense product). It is clearly seen that, the binary systems exhibit a progressively increased resistance to oxidation at high temperature as the HfSi2 is augmented, as compared to monolithic HfB2. Specifically, Figure 7 shows that the normal gain weight remains relatively lower (<0.5 mg/cm2) for all HfB2-based systems exposed to the oxidizing environment when the temperature achieved during the nonisothermal stage is below °C . However , 800 as the temperature is progressively augmented, the situation becomes quite different when the monolithic and composite systems are compared. In particular, th e s amp l e o r i g in a l l y con s i s t ing o f add i t i v e f r e e H fB 2 significantly increased its weight, mostly during the isothermal stage at 1450 °C, thus reaching a value of about 8 mg/cm2 at the end of the experiment. In addition, the curve slope is quite steep, thus indicating the progress of the oxidation process −15 inside the bulk of the sample. In contrast, when the HfB2 vol % HfSi2 specimen was subjected to the same oxidative conditions, the final normal weight gain was markedly lower,  3.4. Oxidation Behavior. The  Article  Figure 7. Specific weight change of dense HfB2 [adapted from Musa et al.6] and HfB2 −xHfSi2 samples (x = 10 and 15 vol %) during TGA oxidation (isothermal run at 1450 °C) performed in air as a function of time.  that is, less than one-fourth of that achieved by monolithic system. Moreover, the TGA curve tends to reach a plateau as a manifestation of the fact that the oxidation is stable at this temperature. The relatively low and thermally stable oxidation rate °C is displayed by the composite materials up to 1450 consistent with the experimental findings observed in the literature when the influence of various Si-containing additives on hafnium diboride is investigated.4−8 More specifically, the behavior observed in Figure 7 can be associated to the corresponding compositional changes of the different systems undergoing oxidation. To this aim, the surface of products obtained at the end of the isothermal TGA tests was analyzed by XRD and the related results are reported in −10 vol % Figure 8 panels a and b for the monolithic and HfB2 HfSi2 composite ceramics, respectively. As far as the first system is concerned, HfO2 was the only phase detected by XRD. This result is a consequence of the fact that HfB2 first reacts with O2 to form HfO2 and B2O3. At relatively low temperature levels, liquid B2O3 (melting point of about 450 °C) fills porous HfO2, thus acting as an diffusion barrier for oxygen. Nevertheless, as the temperature is progressively augmented, B2O3 tends to volatilize due to its high vapor pressure, so that only porous HfO2 is left on the sample surface.6 Consequently, oxygen is allowed to diffuse through the bulk of the material. On the other hand, when the composite system was oxidized, an add i t iona l c ry s ta l l ine s i l ica te pha se , name ly H fS iO4 (hafnon), is also formed (cf. Figures 8b). The presence of th is s i l icate compound is in agreement w ith prev ious investigations related to the high-temperature oxidation of SiC containing HfB2-based composite systems, namely HfB2 −26 20 vol % SiC7 and HfB2 vol % SiCw.6 In this regard, it should be also mentioned that the formation of HfSiO4 at the HfO2/SiC interface during oxidation tests was observed to the material.23 improve the oxidation resistance of A BSE SEM micrograph of the sample cross section relative to the HfB2 −10 vol % HfSi2 system after TGA is shown in Figure 9 along with the corresponding EDS analysis results.  −  9106  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', 'Industrial & Engineering Chemistry Research  Article  Figure 8. Compositional changes of (a) dense HfB2 [adapted from −10 vol % HfSi2 Musa et al.6] and (b) HfB2 after TGA oxidation run at 1450 °C) performed in air. (isothermal  The formation of an oxide layer on the sample surface is evidenced. In addition, as indicated by the EDS analysis (cf. Spectrum 1 of Figure 9), this layer consists of hafnium oxides and silicates. Although the EDS probe used in this work was not able to detect boron in the sample, the latter element is likely incorporated in some silicates phases, according to previous studies reported in the literature on Si-containing ceramics.4−8 HfB2-based The formation of such silicates protects the bulk material from its high temperature oxidation. The EDS analysis corresponding to Spectrum 2 of Figure 9 confirmed the drastic decrease of the oxygen content in the material bulk. Thus, it is possible to conclude that HfSi2 plays a beneficial role for protecting the material from its oxidation. Such behavior is consistent with the experimental findings observed in the literature when investigating the influence of Sicontaining additives to hafnium diboride.4−8  4. SUMMARY AND CONCLUDING REMARKS  (1 − x)HfB2 −xHfSi2 The preparation of (x from 0 to 1) ceramics by self-propagating high-temperature synthesis (SHS) starting from Hf, B, and Si reactants was addressed for the first time in this work. All the systems displayed an SHS character after local ignition, and the generated reaction fronts propagate with velocity and combustion temperature decreasing progressively with the increase of the silicide content. In addition, the complete conversion of reactants into the desired ceramic phases is obtained in all composition taken into account, regardless the HfSi2 percentage. The obtained SHS powders were subsequently consolidated by spark plasma sintering and the influence of the HfSi2 phase as sintering aid was systematically investigated in the range of 0−15 vol %. First, it is observed that the relatively higher densification levels are obtained when starting from composite powders synthesized in a single stage by SHS instead of using mixtures obtained by combining HfB2 and HfSi2 prepared separately.  −10 vol % HfSi2 SPS sample Figure 9. BSE SEM image of the HfB2 after TGA oxidation and related EDS analyses.  Moreover, the minimum HfSi2 content needed to achieve the complete densification of the composite powders under the SPS conditions investigated in the present work (about 1800 °C, P = 50 MPa) is 15 vol %. A 95% dense material can be also obtained using a lower amount of HfSi2 (10 vol %) when the external die diameter was decreased from 35 to 30 mm. As far as the oxidative behavior of the composite sintered material is concerned, TGA experiments performed in air flow up to 1450 °C clearly indicate that the introduction of HfSi2 leads to products with relatively lower and more stable oxidation rate in the temperature range investigated. This behavior is associated to the formation of HfSiO4 and other noncrystalline silicate phases able to hinder the diffusion of oxygen through the bulk of the material. In addition to the considerations above, it is possible to conclude that the use of SHS products is usually more convenient, in comparison to powders obtained by classical methods, as relatively milder sintering conditions are required for the acquisition of bulk materials. This fact represents an important target achieved, particularly when high melting point ceramics have to be consolidated. In addition, from the economical point of view, SHS is an efficient process because of its relatively short synthesis time  9107  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c', 'Industrial & Engineering Chemistry Research  and the occurrence of self-heating to high temperatures instead of external heating as in the classical furnace methods. However, an important aspect to deal with for the practical exploitation of the SHS technology is represented by the maximum amount of powder that can be correspondingly synthesized. Such limit strictly depends on the adopted reactor configuration and size. In this regards, although specific (screw and roll) reactors have been developed for continuous powder production, various practical problems (severe erosion of the reactor, stabilization of the reaction front, difficult-to-control process, etc.) were correspondingly encountered.24 On the other hand, this drawback can be overcome when taking advantage of batteries of batch reactors.24 In this regard, it is worth mentioning that thousands of tons per year of titanium carbide are currently produced following this approach. Of course, also in this case, process control still represents a crucial feature due to the self-propagating character of the synthesis reaction. Details on this subject can be found elsewhere.24 As far as the SPS technology is concerned, the direct passage of the electric pulsed current through the sintering powders and the die containing them leads to very high heating rates, so that processing times can be significantly shortened and sintering temperatures generally lowered, with respect to traditional HP.  ■ AUTHOR INFORMATION  Corresponding Author  *E-mail: roberto.orru@dimcm.unica.it. Tel.: +39-070-6755076. Fax: +39-070-6755057.  Notes  The authors declare no competing financial  ■ ACKNOWLEDGMENTS  interest.  The financial support from MIUR (Italy) through the SUPERSOLAR project (Prot. RBFR12TIT1) FIRB 2012 “Futuro in Ricerca” is gratefully acknowledged. The authors also thank Dr. Paola Meloni (University of Cagliari, Italy) and Dr. Eng. Gianfranco Carcangiu (National Research Council, Cagliari, Italy), for their valuable support in SEM investigation. G.C. recalls with fond memory the work conducted under the direction of his mentor, professor Massimo Morbidelli.  ■ REFERENCES  J. A. Soc.  (1) Fahrenholtz, W. G.; Hilmas, G. E.; Talmy, I. G.; Zaykoski, Refractory diborides of zirconium and hafnium. J. Am. Ceram. 2007, 90, 1347. (2) Cheminant-Coatanlem, P.; Boulanger, L.; Deschanels, X.; Thorel, A. Microstructure and nanohardness of hafnium diboride after ion irradiations. J. Nucl. Mater. 1998, 256, 180. (3) Orru ̀, R.; Licheri, R.; Locci, A. M.; Cincotti, A.; Cao, G. Consolidation/synthesis of materials by electric current activated/ assisted sintering. Mater. Sci. Eng. R 2009, 63, 127. R.; Orru ̀, (4) Licheri, R.; Musa, C.; Locci, A. M.; Cao, G. Consolidation via spark plasma sintering of HfB2/SiC and HfB2/ HfC/SiC composite powders obtained by self-propagating highJ. Alloys Compd. 2009, 478, 572. temperature synthesis. (5) Sciti, D.; Bonnefont, G.; Fantozzi, G.; Silvestroni, L. Spark plasma sintering of HfB2 with low additions of silicides of molybdenum and tantalum. J. Eur. Ceram. Soc. 2010, 30, 3253. (6) Musa, C.; Orru ̀, R.; Sciti, D.; Silvestroni, L.; Cao, G. Synthesis, consolidation and characterization of monolithic and SiC whiskers reinforced HfB2 ceramics. J. Eur. Ceram. Soc. 2013, 33, 603. (7) Hinze, J. W.; Tripp, W. C.; Graham, H. C. The high-temperature oxidation behaviour of a HfB2 + 20 v/o SiC composite. J. Electrochem. Soc. 1975, 122, 1249.  Article  an  (8) Monteverde, F.; Bellosi, A. The resistance to oxidation of −SiC composite. HfB2 J. Eur. Ceram. Soc. 2005, 25, 1025. (9) Kuznetsova, S. V.; Kuznetsov, S. A. Electrochemical synthesis of hafnium silicides. Russ. J. Electrochem. 2009, 45, 749. (10) Yen, B. K.; Aizawa, T.; Kihara, J.; Sakakibara, N. Reaction synthesis of refractory disilicides by mechanical alloying and shock reactive synthesis techniques. Mater. Sci. Eng., A 1997, 239−240, 515. (11) Monteverde, F. Hot pressing of hafnium diboride aided by different sinter additives. J. Mater. Sci. 2008, 43, 1002. (12) Bronson, A.; Ma, Y. T.; Mutso, Rein R. Compatibility of refractory metal boride/oxide composites at ultra-high temperatures. J. Electrochem. Soc. 1992, 139, 3183. (13) Park, J.-H.; Yoon, J.-K.; Doh, J.-M.; Lee, B.-S.; Shon, I.-J. Simultaneous high-frequency induction heated combustion synthesis and consolidation of nanostructured HfSi2-SiC composite. Ceram. Int. 2009, 35, 1677. (14) Mun ir , Z . A . ; Anse lm i-Tambur in i , U . Se l f-propagat ing exothermic reactions: The synthesis of high-temperature materials by combustion. Mater. Sci. Rep. 1989, 3, 277. (15) Licheri, R.; Orru ̀, R.; Locci, A. M.; Cao, G. Efficient synthesis/ sintering routes to obtain fully dense ZrB2-SiC ultra-high-temperature ceramics (UHTCs). Ind. Eng. Chem. Res. 2007, 46, 9087. (16) Licheri, R.; Orru ̀, R.; Musa, C.; Cao, G. Combination of SHS and SPS techniques for fabrication of fully dense ZrB2-ZrC-SiC composites. Mater. Lett. 2008, 62, 432. (17) Licheri, R.; Orru ̀, R.; Musa, C.; Cao, G. Efficient for the fabrication of dense TaB2-based ultra high ceramics. ACS Appl. Mater. Interfaces 2010, 2, 2206. (18) Cincotti, A.; Licheri, R.; Locci, A. M.; Orru ̀, R.; Cao, G. A review on combustion synthesis of novel materials: Recent experimental and J. Chem. Technol. Biot. 2003, 78 (2−3), 122. modeling results. (19) Locci, A. M.; Orru ̀, R.; Cao, G.; Munir, Z. A. Simultaneous spark and densification of TiC−TiB2 plasma synthesis composites. J. Am. Ceram. Soc. 2006, 89, 848. (20) Barin, I. Thermochemical Birmingham, UK, 1989. (21) Mishra, S. K.; Das, S.; Pathak, L. C. Defect structures in zirconium diboride powder prepared by self-propagating hightemperature synthesis. Mater. Sci. Eng., A 2004, 364, 249. (22) Zhao, J.-C.; Bewlay, B. P.; Jackson, M. R.; Chen, Q. Hf−Si binary phase diagram determination and thermodynamic modelling. J. Phase Equilib. 2000, 21, 40. (23) Mergia, K.; Liedtke, V.; Speliotis, Th.; Apostolopoulos, G.; Messoloras, S. Thermo-mechanical behaviour of HfO2 coatings for aerospace applications. Adv. Mater. Res. 2009, 59, 87. (24) Cincotti, A.; Orru ̀, R.; Pisu, M.; Cao, G. Self-propagating reactions for environmental protection: Reactor engineering aspects. Ind. Eng. Chem. Res. 2001, 21, 5291.  technologies temperature  Data  Substances;  VHC:  of  Pure  9108  dx.doi.org/10.1021/ie4032692 |  Ind. Eng. Chem. Res. 2014, 53, 9101−9108  \\x0c']"
},{
  "_id": 257,
  "PDF": "TEM analysis, mechanical characterization and oxidation resistance of a highly refractory ZrB2 composite.pdf",
  "Text": "['Journal of Alloys and Compounds 602 (2014) 346-355  Contents lists available at ScienceDirect  Journal of Alloys and Compounds  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / j a l c o m  TEM analysis, mechanical characterization and oxidation resistance of a highly refractory ZrB2 composite  ⇑  Laura Silvestroni  , Diletta Sciti  CNR-ISTEC,  Institute of Science and Technology for Ceramics, Via Granarolo 64,  I-48018 Faenza,  Italy  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 20 January 2014 Received in revised form 20 February 2014 Accepted 21 February 2014 Available online 28 February 2014  Keywords:  ZrB2 Silicides TEM High temperature mechanical properties Oxidation  The microstructure of a dense ultra-high-temperature ceramic, namely ZrB2 containing 15 vol% of WSi2, was characterized by X-ray diffraction, scanning and transmission electron microscopy. ZrB2 displayed a core-shell structure: the core was constituted by the original MB2 grain and the shell by a (Zr,W)B2 solid solution which grew epitaxially on the core. According to the final microstructure with mixed Zr,W-silicides, -borides and -carbides at the triple points and clean grain boundaries, a densification mechanism for ZrB2 in presence of WSi2 is proposed. This composite displayed excellent thermomechanical properties, like strength retention of 84% at 1500 °C in air and moderate oxidation up to 1650 °C. Correlation between microstructure and properties are here presented in relationship to other ultra-refractory ceramics available in the literature.  Ó 2014 Elsevier B.V. All rights reserved.  1. Introduction  The impelling demand of reusable and long lasting materials, able to operate at high temperature and for long time pushes the scientific research towards continuous search of materials possessing a combination of properties always more challenging. The new concept of hypersonic vehicles actually lies on the availability of these new materials. Like most of parts of motors, the components of propulsion or hypersonic vehicles are exposed to significant heating flux where mechanical solicitations are present too. Therefore the main requirement that a material for such applications should satisfy is a combination of refractoriness at high temperature in oxidizing environment and sufficient mechanical resistance in such environment. The class of compounds commonly known as ultra-high temperature ceramics (UHTCs) has been identified as potential candidate for operating in harsh conditions owing to their melting point exceeding 3000 °C and their resistance to ablation [1]. Among the most investigated systems, ZrB2-based composites certainly hold supremacy in literature and overall knowledge of the microstructure-properties relationships over the other borides and carbides compounds. It has been shown that strength up to 1 GPa [2] can be achieved upon a tailored design of composition and processing. A recent work reported that texturing of ZrB2 by  ⇑ Corresponding author. Tel.: +39 546 699723; fax: +39 546 46381. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  http://dx.doi.org/10.1016/j.jallcom.2014.02.133 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.  magnetic field during slip casting process enabled to achieve strength overcoming 800 MPa at 1600 °C [3]. It has further been demonstrated that the addition of dopants, in particular silicides and carbides of transition metals in groups V and VI, can induce a strength retention at high temperature by altering the microstructural features [4,5]. For example, strength of 640 MPa at 1500 °C in argon has been reported after introduction of WC [5]. In another paper, WC-doping was introduced in ZrB2ASiC composite in amounts ranging from 3 to 7 vol% and again increased oxidation performances of the W-doped ceramics as compared to pure ZrB2ASiC were reported [6]. The reasons for this improvement were ascribed to the formation of a eutectic between WO3 and ZrO2 at 1275 °C, that acted as liquid phase sintering and decreased the porosity of the scale, or WO3 acted as barrier itself due to the volume increase associated with oxidation of W to WO3. In a dedicated study, the oxidation behavior of ZrB2 with 0, 4, 6, or 8 mol% W was studied at 800-1600 °C and pointed out that the addition of W into B2O3 increased the stability of the protective glassy layer, which resulted in higher oxidation resistance [7]. Other authors suggested that tungsten oxide species form strong acid sites in ZrO2 and inhibit ZrO2 tetragonal to monoclinic structural transformations [8]. In this work, a ZrB2 ceramic was hot pressed through addition of WSi2 which enabled the densification and the development of a microstructure possessing high mechanical strength at 1500 °C in air and relatively low oxidation rate up to 1650 °C. WSi2, with a tetragonal structure, has been the focus of considerable attention as an attractive material for electronics and high temperature  \\x0c', 'L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  347  application, as its properties provide the desirable combination of high melting temperature (2160 °C), good strength at high temperature, good creep resistance and high oxidation resistance which relies on the formation of vitreous, dense, adherent SiO2 scale [9]. Following previous positive results obtained with other transition metal silicides [10] and in view of the above mentioned physico-chemical features, WSi2 could be an adequate sintering additive for ZrB2 for applications in ultra-high temperature environments. Different analytical techniques were used in order to elucidate the composition and structure of the newly formed W-based phases. Besides the conventional characterization by X-ray diffraction and scanning electron microscope, transmission electron microscope was here used as main tool to explore microstructures on a small length scale to give information about the densification mechanisms and to establish correlations between microstructural features and bulk material properties.  2. Experimental procedure  A composite containing 15 vol% of sintering additive was prepared using the following commercial powders: hexagonal ZrB2, Grade B (H.C. Starck, Germany), specific surface area 1.0 m2/g, impurity max content: C: 0.25 wt%, O: 2 wt%, N: 0.25 wt%, Fe: 0.1 wt%, Hf: 0.2 wt%, particle size range 0.1-8 lm; tetragonal WSi2, Italy), \\x00325 mesh, 99.5%, traces of metals <6000 ppm. (Sigma Aldrich, Milano, The powder mixture was ball milled with zirconia media in a polyethylene jar for 24 h in absolute ethanol at 130 rpm. The ratio between powder, solvent and milling media was 1:1:1 in weight. Subsequently the slurry was dried in a rotary evaporator. Sintering was conducted in a hot pressing in low vacuum (\\x18100 Pa) using an induction-heated graphite die with a constant uniaxial pressure of 30 MPa, heating rate 20 °C/min and free cooling. The maximum sintering temperature was 1930 °C. The bulk density was measured by Archimedes’ method. To identify the crystalline phases formed, the sample was examined using X-ray diffraction (Siemens D500, Karlsruhe, Germany), with Cu Ka radiation, step size of 0.04 and 1 s counting rate. The microstructure was analyzed on fractured and polished surfaces by scanning electron microscopy (FE-SEM, Carl Zeiss Sigma NTS Gmbh Öberkochen, Germany) and energy dispersive X-ray spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK). TEM samples were prepared by cutting 3 mm discs from the bulk of the sintered pellets. These were mechanically ground down to about 20 lm and then further ion beam thinned until small perforations were observed by optical microscope. Local phase analysis was performed using transmission electron microscopy (TEM) equipped with an energy-dispersive X-ray system. TEM analysis was performed by a FEI Tecnai F20 ST, with an acceleration voltage of 200 kV, equipped with an EDAX EDS X-ray spectrometer PV9761 with super ultra-thin window. Quantitative calculations of the microstructural features, such as residual porosity and secondary phases content, were carried out via image analysis with a commercial software package (Image Pro-plus 4.5.1. Media Cybernetics, Silver Spring, MD, USA). Vickers microhardness Zwick 3212 tester. Fracture toughness (KIc) was evaluated using chevron-notched beams (CNB) in flexure. The test bars, 25 \\x02 2 \\x02 2.5 mm3 (length by width by thickness, respectively), were notched with a 0.1 mm-thick diamond saw; the chevron-notch tip depth and average side length were about 0.12 and 0.80 of the bar thickness, respectively. The specimens were fractured using a semi-articulated silicon carbide four-point fixture with a lower span of 20 mm and an upper span of 10 mm using a screw-driven load frame (Instron mod. 6025). The specimens were loaded with a crosshead speed of 0.05 mm/min. The ‘‘slice model’’ equation of Munz et al. [11] was used to calculate KIc. On the same machine and with the same fixture, the flexural strength (r) was measured at room temperature and at 1500 °C in air on chamfered bars 25 \\x02 2.5 \\x02 2 mm3 (length \\x02 width \\x02 thickness, respectively), in 4-point bending using a crosshead speed of 0.5 mm/min. Before the bending test, a soaking time of 18 min was set to reach thermal equilibrium. 5 specimens were tested at room temperature and 3 at high temperature. The oxidation resistance was tested at 1200, 1350, 1500 and 1650 °C on 13 \\x02 2.5 \\x02 2 mm3 bars in static air in a bottom loading furnace box (Nannetti FC/ 18, Faenza, Italy). The specimens were located in the furnace on ZrO2 supports when the maximum temperature was reached and then removed and air quenched after an exposure time of 15 min. All the specimens were previously cleaned in acetone. The mass of the specimens was measured before and after exposure. The microstructural modifications induced in the oxidized specimens were evaluated by XRD and SEM-EDS on surface and cross section.  (HV1.0) was measured with a load of 9.81 N using a  3. Results  3.1. Densification behavior  The densification curve recorded during hot pressing run is depicted in Fig. 1a where density and temperature are reported as a function of time. It can be read that significant shrinkage started just above 1750 °C (TON), which is a relatively high temperature as compared to other transition metal silicides (ZrSi2: 1350 °C, MoSi2: 1500 °C [10]). The inset in Fig. 1a shows the maximum densification rate for different ZrB2 composites containing the same  Fig. 1. (a) Densification curve recorded during hot pressing. TON: temperature at which shrinkage begins, TLP: temperature at which liquid phase forms. In the inset: plot of the maximum densification rate as a function of the temperature for different silicides; Z: ZrSi2, T: TaSi2, M: MoSi2, W: WSi2. (b) X-ray diffraction pattern of the sintered composite showing the crystalline phases and (c) ZrB2 peaks splitting at high 2-Theta angles.  \\x0c', '348  L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  amount of various silicides, i.e. ZrSi2, MoSi2, TaSi2 and WSi2 [10]. It is evident that WSi2, besides favoring the sintering process at the highest temperature, also enables the densification at the lowest speed. The sintering process proceeded at the same rate up to 1850 °C, after which a steeper slope indicated the improvement of densification mechanisms, possibly assisted by liquid phases (TLP). At 1930 °C the relative density was around 87% of the theoretical value, but a dwell at this temperature for 25 min enabled the achievement of the full density.  3.2. Microstructural  features of the as sintered ceramic  The X-ray diffraction pattern of the composite is reported in Fig. 1b; a Si standard was placed on the top of the specimen in order to detect any peak shift. In the whole 2-Theta range, ZrB2 was the main crystalline phases, but reflections of tetragonal WB were found too, whilst WSi2 was present just in traces. At high 2-Theta angles, above 80° (Fig. 1c), another phase is clearly visible on the right of ZrB2 (PDF# 34-0423). The unit cell parameters of this newly formed phase were a = 3.1653 and c = 3.5244 Å, i.e. slightly shorter than those of pure ZrB2 (a = 3.1690 and c = 3.5300 Å), which indicates a contraction of the unit cell. The presence of these additional reflections was interpreted as due to the formation of a (Zr,W)B2 solid solution. However, the exact composition of this new phase was not computable through Vegard’s rule, owing to the too small amount of W entered into ZrB2 lattice, within the instrumental error. Images of the microstructure by SEM are shown in Fig. 2. Fig. 2a shows an overall view of the polished section: the grains are rounded with average dimensions of 3.1 ± 0.8 lm and are homogeneously surrounded by a brighter gray phase. Different gray levels in the matrix are due to different grains orientation. The main secondary phases are about 2 vol% of silica, with dark contrast, and 3.5% of a white phase containing W,B,C,O. Small SiC grains were  often found next to silica pokets. No major defects or porosity were observed. Fig. 2b discloses the morphology of the matrix: the core is ZrB2, whilst the outer brighter region, the shell, is ZrB2 containing 2 at.% of W, resulting in a solid solution like (Zr0.98W0.02)B2, as the EDS spectra and elemental mapping in Fig. 2c and d demonstrate. The presence of carbon in the spectra is due to beam contamination. Images acquired with in-lens electrons, Fig. 2e-g, revealed that the bright phase observed by backscattered electrons in Fig. 2a is actually composed by several other phases: the largest areas are WCO and WBC, always adjacent to each other’s and with low dihedral angles suggesting solidification from liquid phase. At the extremities, brighter regions with generally elongated shape are WB with Zr traces, Fig. 2e and f, whilst the bright phases found at the triple points with concave shape are mixed Zr,W-silicides with stoichiometry difficult to identify owing to the superimposition of W and Si lines, Fig. 2g. TEM investigations enabled a more precise definition of the constituent phases. Examples of the matrix grain morphology are reported in Fig. 3a-c. A peculiar feature of the diboride grains was the presence of dislocations in the shell. Similarly to previous study on analogous composites [10,12], core and shell were epitaxial, as disclosed by electron diffraction patterns, Fig. 3. In spite of that, no modification of lattice parameters was detected by SAEDP technique, owing to the small differences in atomic radii between Zr and W (0.160 and 0.137 nm, respectively) and to the little W content in the shell. The boundary between core and shell was quite sharp and defined by dislocations pile up, as depicted in Fig. 3a-c. W-rich precipitates were occasionally found at the core-shell interface, as the inset in Fig. 3c shows. TEM EDS quantitative analyses performed on core and shell, confirmed around 2 wt% of W in the solid solution, Fig. 3. The study of secondary phases and triple points by electron diffraction and EDS confirmed the formation of large WB grains, with WC or WSi2 smaller grains at the apical parts, Fig. 4a-e and the  Fig. 2. SEM images of the sintered ceramic showing (a) the overall microstructure, (b) the matrix morphology with (c) EDX spectra of core and shell. (d) Elemental mapping of the matrix showing the distribution of W in the shell. (e)-(g) secondary W-based phases.  \\x0c', 'L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  349  Fig. 3. (a)-(c) TEM images showing the core-shell structure of the matrix of the sintered ceramic with EDX spectra at low and high voltage and SAEDP taken from (a) confirming epitaxy between ZrB2 and (Zr,W)B2. The inset in (c) shows a W-rich precipitate occasionally found at the core-shell interface.  Fig. 4. TEM images of the secondary phases with EDX spectra and diffraction patterns showing the presence of WB, WC, WSi2 and residual ZrO2.  \\x0c', '350  L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  segregation of ZrO2 or impurities, like Fe-rich phases, at the triple junctions, Fig. 4f. Analyses of the interfaces revealed non wetted grain boundaries between adjacent (Zr,W)B2 grains, Fig. 5a-f, and between WB and (Zr,W)B2, Fig. 5g-i. Clean interfaces were found both in case of simple triple junctions, Fig. 5a-c, and in case of segregation of secondary phases at the triple points, Fig. 5d-f. Strain contrasts were generally observed in ZrB2 grains when other phases crystallized at the junctions, Fig. 4a and e. To obtain statistically valid results, several interfaces were analyzed and no amorphous film was found between (Zr,W)B2, (Zr,W)C, (W,Zr)Si2 and WB grains, as confirmed in Fig. 6.  3.3. Mechanical properties  The mechanical properties measured on the ZrB2AWSi2 composite are reported in Table 1 together with properties of other ZrB2 ceramics sintered with addition of silicides (MoSi2, TaSi2) or W-compounds (WC) [5,10,13,14]. As expected, the fracture toughness of this composite was in the range of other ZrB2-based ceramics measured with the same chevron v-notched method, 3.6 MPa m1/2 [10]. The hardness, around 18 GPa, resulted higher than other ZrB2ceramics sintered with addition of other transition metal silicides [10], despite coarser grain size and 2 vol% of silica, probably thanks  to the presence of other hard secondary phases, like WB and WC. The room temperature flexural strength was 640 with very low standard deviation, less than 3%, confirming a uniform microstructure without large flaws. Interesting results were obtained for flexural strength tests at 1500 °C in air: a retained strength of more than 84% was preserved, outdoing most of ZrB2 compounds available in the literature. Excellent results were obtained also by Zou and coworkers [5] who even measured a strength increase, however those data are not directly comparable with those of the present work, since their tests were performed in argon and in 3-point bending.  3.4. Microstructural  features of the oxidized ceramic  When we deal with oxidation of dark ceramics like borides, an increase of the oxidation temperature induces the modification of the gray color, typical of the as sintered material, to dark gray, index of glass formation, to whitish-yellow, index of formation of the corresponding metal oxide. In the present case, the oxidized specimens remained of gray color up to 1350 °C, but when the ceramic was oxidized above 1500 °C it finished with whitish-yellowish aspect. The main crystalline phase detected by X-ray diffraction on the oxidized samples, was monoclinic ZrO2. In addition tungsten oxide was detected at all temperatures, Fig. 7. Traces of WB were instead  Fig. 5. TEM and HRTEM images of the triple point junctions showing clean grain boundaries between (a)-(f) adjacent (Zr,W)B2 grains and (g)-(i) between WB and (Zr,W)B2.  \\x0c', 'L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  351  Fig. 6. HR-TEM images showing clean grain boundaries between (Zr,W)B2, (Zr,W)C, (W,Zr)Si2 and WB.  Table 1 Composition, mean grain size and mechanical properties of ZrB2AWSi2 ceramic compared to other ZrB2-composites containing either transition metal silicides or W-compounds. m.g.s.: mean grain size, KIc: fracture toughness measured by chevron notch technique, HV: microhardness (1 kg), r: 4-point flexural strength at room temperature or at 1500 °C in air.  Composition (vol%)  ZrB2 + 15 WSi2 ZrB2 + 15 MoSi2 ZrB2 + 15 TaSi2 ZrB2 + 3 WC (PLS) ZrB2 + 3 WC (HP) ZrB2 + 20SiC + 5 WC (ZrB2 + 20SiC) + 10 WC  a  SEVNB. b 3-pt Beniding strength. c Ar.  ZrB2 m.g.s. (lm)  KIc (MPa m0.5)  3.1 ± 0.8 2.4 2.0 9.1 ± 5.6 4.0 ± 1.2 1.1 -  3.62 + 0.35 3.50 ± 0.60 3.80 ± 0.10 - - - 6.5 ± 0.2a  HV (GPa)  18.3 + 0.4 14.9 ± 0.5 17.8 ± 0.5 14.5 ± 2.6 23.0 ± 0.9 - 13.9 ± 0.2  rRT (MPa)  r1500 (MPa)  641 + 19 704 ± 98 840 ± 33 444 ± 30 565 ± 53 605b 518 ± 10b  537 + 16 333 ± 31 374 ± 5 - - 640b,c -  Ref.  This work [10] [10] [13] [13] [5] [14]  Fig. 7. X-ray diffraction spectra of ZBW oxidized at 1200, 1350, 1500, 1650 °C for 15 min.  \\x0c', '352  L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  visible up to 1350 °C, whilst ZrB2 reflections were not identified at any oxidation temperature. The weight gain per unit surface area (not shown) increased with the temperature, moderately up to 1350 °C, 3 mg/cm2, then more at higher temperature, i.e. up to 12 mg/cm2 at 1650 °C. Images of the external surface at the various oxidation temperatures are given in Fig. 8a-d, whilst pictures of the fractured cross sections are reported in Fig. 8e-h with EDS spectra of the mentioned phases. As first observation, it can be seen that the thickness of the oxidized layer progressively increased with the temperature and passed from 24 lm after 1200 °C, to 55 lm at 1350 °C, to 135 lm at 1500 °C and to 210 lm at 1650 °C. At 1200 °C, Fig. 8a, the surface was characterized by large areas of glassy B2O3 and submicrometric grains of ZrO2 which were covered by tiny WO3 whiskers, inset in Fig. 8a, bottom. Occasionally this phase was also found in form of agglomerates with bright contrast among ZrO2 grains, see the inset in Fig. 8a, top. In the cross section, Fig. 8e, about 24 lm underwent modification and this scale was mostly constituted by ZrO2 grains with intergranular WO3 agglomerates adherently attached to the unreacted bulk without visible cracks and topped with discontinuous B2O3 glassy layer. At 1350 °C boron oxide disappeared from the surface and, in the outermost layer, only ZrO2 grains immersed in silica glass were detected, Fig. 8b. Even W traces were barely detected among the crystals. The scale thickness doubled as compared to the sample oxidized at 1200 °C, but maintained the same morphology, Fig. 8f. Also in this case the oxidized layer was well adherent to the matrix, inset in Fig. 8f. The oxidation at 1500 °C induced a notable variation of the external morphology which appeared wavy and rutty showing  bumps and craters. The ZrO2 grains coarsened to about 2 lm and appeared aggregated and melted on a silica continuous layer, Fig. 8c. The growing planes were well visible in ZrO2 and W-rich phases with bright contrast were found at the grain boundary, Fig. 8c. Differences in coefficient of thermal expansion between zirconia and silica provoked the formation of microcracks. This oxidation induced a diversification in morphologies of the ZrO2 sub layer, Fig. 8g. Starting from the top, 20 lm coarse ZrO2 grains were standing on another 20 lm of columnar ZrO2, which was topping about 95 lm of dense rounded-grains ZrO2. The whole thickness was partially filled with silica and the ZrO2 grains were coated with homogeneously dispersed WO3 drops which eventually turned into 5 lm big agglomerates, see the insets in Fig. 8g. The surface of the sample oxidized at 1650 °C was not notably different from the one oxidized at 1500 °C, just more bumps, craters and ‘‘cabbage-like’’ structures were noticed, Fig. 8d. Similarly, the morphology of the various scales remained basically unmodified as compared to the features observed for the samples oxidized at 1500 °C, Fig. 8h, but each scale just increased in thickness and columnar ZrO2 layer further developed and displayed microcracks within the elongated grains.  4. Discussion  4.1. Microstructure evolution during densification  In the sintered material, a notable change of the microstructural features occurred, which is strictly related to the densification mechanisms activated by WSi2. Addition of metal disilicides, MeSi2 (Me@Ti, Cr, Zr, Mo, Ta, W), to TiB2 and ZrB2 was first investigated by Pastor and Meyer in 1974 [15] and, according to their studies,  Fig. 8. SEM images of (a)-(d) external surface and (e)-(h) cross section of the oxidized ceramic. From left to right: 1200, 1350, 1500, 1650 °C with microstructural details in the insets. In the bottom: EDS spectra of B2O3 glass in (a), ZrB2 and ZrO2 in (f), WO3 in (g) and SiO2 from all surfaces.  \\x0c', 'L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  353  densification of borides with MeSi2 was successful because it was conducted at temperatures close to the MeSi2 melting point. For WSi2, however, full densification is achieved although the temperature range is well below its melting point, 2160 °C [16]. However, low dihedral angles suggest WSi2 was ductile at the sintering temperature. According to the phase diagrams involving WASi [16], WAB [17], WAC [17], no liquid phases are predicted to form below 1930 °C. Also the eutectic at 1275 °C [18] between the oxides covering the particles, WO3 and ZrO2, does not seem to affect the beginning of the shrinkage occurring at 1750 °C, TON in Fig. 1a. Only Si, SiO2 and B2O3 are liquid below the sintering temperature. However, the presence of W-based phases at the triple junctions containing Zr traces suggests the existence of some liquid phases in which ZrB2 is partially soluble in. Similar to MoSi2, and TaSi2 [10,12], the densification of ZrB2AWSi2 composite is presumed to be assisted by a transient liquid phase, due to the reaction between WSi2 and the oxide impurities covering the diboride powders, according to reaction (1):  2 WSi2 þ B2O3 ðlÞ ! 2 WB þ 2:5 SiðlÞ þ 1:5 SiO2 ðlÞ  ð1Þ  Reaction (1) implies removal of B2O3 from the surface of the diboride particles, formation of SiA and SiAO-based liquids and range, \\x0048 kJ/ is favorable in the whole sintering temperature mol at 1700 °C. Reaction (1) is not exhaustive and, in the reducing environment of the hot pressing chamber, it is likely that variants of the same reaction locally occur, for example resulting in the formation of WO3, reaction (2):  9 WSi2 þ 4 B2O3 þ 9 COðgÞ ! 8 WB þ WO3 þ 9 SiC þ 9 SiO2 ðlÞ  ð2Þ  TLP  which has a negative Gibbs free energy as well, \\x0036 kJ/mol at 1700 °C. Nevertheless, all these intermediate phases concur to the formation of liquid phase at temperature around 1850 °C, in Fig. 1a, that promotes mass transfer mechanisms, by partial dissolution of the boride matrix. A distinctive feature of the ZrB2AWSi2 composites is the formation of (Zr,W)B2 solid solution around ZrB2 grains, which was already found for other ZrB2AMeSi2 and it was also assessed that (Zr,Me)B2 shells grow epitaxially onto ZrB2 cores [10,12]. This grains morphology suggests again a solution re-precipitation mechanism, summarized by reaction (3), even if diffusion of W into the diboride lattice and Zr into the silicide cannot be ruled out, as previously observed for analogous systems involving different transition metal cations which are close in the periodic table [19]:  ZrB2 þ WSi2 ! ðZr; WÞB2 þ ðW; ZrÞSi2  ð3Þ  During cooling the transient liquid-phase solidifies, resulting in the formation of discrete crystalline phases, such as the observed WB, WC. At the same time, the carbo-reduction of SiO2A and WO-based species to SiC and WC is favorable too:  SiO2 þ C ! SiC þ CO2 ðgÞ  WO3 þ 2C ! WC þ COðgÞ þ CO2 ðgÞ  ð4Þ  ð5Þ  The chemistry of the crystalline phases at the triple point junctions seems to confirm the above reactions to occur. We can state that, compared to previously studied systems, the densification of this one with WSi2 proceeds through more sluggish mechanisms even if the sintering temperatures is higher, see the inset in Fig. 1a. This can be determined by combination of factors: the initial liquid amount is limited and more refractory, its viscosity progressively increases owing to the incorporation of higher and higher amount of cations from boride and silicides,  and also silica glass at such high temperatures tends to be carboreduced to SiC. In addition, the study of the interfaces also reveals that the liquid does not wet the matrix, as all the grain boundaries were clean from any intergranular phases.  4.2. Fracture strength  The featuring characteristic of this composite is its very high refractoriness at 1500 °C in air, as testified by fracture strength of 540 MPa at this temperature, namely 84% of the room temperature value, and the linearity of the load-displacement curve, displayed in Fig. 9a. An example of the fracture surface after bending test at 1500 °C is shown in Fig. 9b. On the cross section, about 150 lm of material underwent oxidation and this layer was formed by ZrO2 rounded grains containing W traces and covered by H3BO3 acicular crystals deriving from hydration of B2O3 after exposure to humidity in the environment. This feature is quite surprising, since B2O3 has low vapor tension and evaporates above 1000 °C. The fact that we still see crystalline boron oxide in the oxide scale could be due to the presence of tungsten oxide which reduced the volatility of the borate glass, as previously stated in a dedicated study [7]. The strength retention at high temperature is usually attributed to the formation of a protective oxide scale [20], crack healing or to the release of internal residual stresses [21]. In this system, neither amorphous phase nor glassy phases were found in the microstructure. Crack healing can be therefore ruled out as strengthening mechanism. Regarding the relaxation of the internal residual stresses, the test temperature of our specimens, 1500 °C, was not close  Fig. 9. (a) Load-displacement curves for flexural strength at room temperature and at 1500 °C in air. (b) SEM images of the cross section of a bar after flexural strength testing at 1500 °C showing the presence of acicular BO-based crystals.  \\x0c', '354  L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  to the material sintering temperature, 1930 °C, so that the relaxation could not have been truly effective. In a recent TEM study conducted upon high temperature flexural strength on ZrB2AB4C composites with clean grain boundaries like in the present case [3], it was pointed out that, below a testing temperature of 1600 °C, no interconnected cavities developed at the grain junctions and that the stress induced by the load was progressively accommodated by ZrB2 grains plastic deformation, due to the generation and flow of dislocations. This mechanism promoted a strength increase at high temperature in that composite [3]. In the ceramic of the present work, we observed dislocation networks already after sintering, as Fig. 3 clearly shows. The absorption of small amount of mechanical energy during the flexural test by dislocation flow and multiplication could give a partial contribution to the remarkable high temperature strength of ZrB2AWSi2 composite. Also the clean grain boundaries and the presence of highly refractory secondary phases at the triple points, like WB, possessing a melting point of 2655 °C [17], or WC, 2870 °C [17], Fig. 4, could be beneficial for the high temperature strength by preventing grain boundary sliding at high temperature. In conclusion, the high refractoriness of the ZrB2AWSi2 ceramic is very likely a combination of multiple factors: the formation of a stable and protective borate glass on the surface during oxidation, the intrinsic refractoriness of the crystalline compounds, the absence of amorphous phases at the grain boundaries and the tendency of the matrix grains to absorb energy by plastic deformation due to dislocations flow.  4.3. Oxidation behavior  From the microstructure evolution presented in Section 3.4, it is obvious that multiple and various phenomena occur upon oxidation of ZrB2 doped with WSi2 implying melting, formation of glassy phases, phase transformation and vaporization. The beneficial effect of tungsten on the oxidation behavior of ZrB2 and ZrB2ASiC composites has been already outlined in previous works [5-7]. The improved behavior of W-doped ceramics was ascribed to different phenomena:  \\x0f The formation of an eutectic between WO3 and ZrO2 at 1275 °C, that acted as liquid phase sintering and decreased the porosity of the outer scale [5,6]. \\x0f The fact that WO3 performed as barrier itself due to the volume increase associated with oxidation of W-species to WO3 [6]. \\x0f The incorporation of W in the borate glass which increased its stability [7].  In the present case, at 1200 °C, boride, carbide and silicide and are expected to oxidize according to:  ZrB2 þ 5=2O2 ! ZrO2 þ B2O3  WSi2 þ 7=2O2 ! WO3 þ 2SiO2  WC þ 2O2 ! WO3 þ CO  2WB þ 9=2O2 ! 2WO3 þ B2O3  ð6Þ  ð7Þ  ð8Þ  ð9Þ  is al In this temperature range, a subsurface oxidation layer ready formed, mainly composed of ZrO2, for around 25 lm. This picture remains essentially the same at 1350 °C. At 1500 °C the ceramic starts to oxidize readily and both weight gain and oxide thickness notably increase, probably owing to the formation of liquid phases. When W-compounds are present, litemperatures lower than 1300 °C, due to quid phase can form at the WO3AZrO2 eutectic [18]. Additional sources of liquid come from mixed (Zr,W)-silicides, oxides and carbides present as  secondary phases in the sintered microstructure. Liquid phases provide a fast path for oxygen transport and enhance oxidation. In addition, the condensation of tungsten oxide, coming from reactions (7)-(9), from vapor phase is well evident, as it is found as nanosized rounded particles on ZrO2 grains, Fig. 8g and h. At 1650 °C, the aforementioned phenomena are further enhanced. Liquid phases cause convection phenomena which make zirconia grow in columnar structures, as consequence of convective transport of low viscosity borosilicate liquid, similar to what observed for ZrB2ASiC composites [22,23]. As a matter of fact, differently from what reported in [6], the formation of solid solutions between tungsten oxide and zirconia was not proved by SEM-EDS, although the amount of available W was more than double than the minimum amount of W reported to be effective in promoting liquid phase sintering of ZrO2 and thus hindering the oxidation [6]. Actually, the oxidation protection imparted by W addition is reported to be better above 1600 °C [6] or even above 1800 °C [24]. Further oxidation studies at higher temperature are in progress to assess the stability of the external scale in this more severe regime, as slightly suggested by the weight gain trend, whose slope tends to decrease.  5. Conclusions  A ZrB2-based ceramic was hot pressed to fully density at 1930 °C through addition of WSi2. Although the high sintering temperature, the microstructure resulted fine and homogeneous, thanks to the beneficial effect of WSi2 and its derivatives products. Core-shell structures were found in the boride matrix and TEM analyses revealed the segregation of highly refractory compounds at the triple points, like WB, WC and WASi phases with clean grain boundaries. These features resulted favorable for the high temperature performances. In particular, the mechanical strength at 1500 °C in air was 540 MPa, about 84% of the room temperature value. Possible reasons explaining this behavior were the incorporation of W in the borate glass during oxidation, which increased the oxide layer stability and protective action, and to the presence of clean grain boundaries and highly refractory secondary phases at the triple junctions. The oxidation behavior was characterized in the 1200-1650 °C temperature in a bottom loading furnace range and this composite underwent modifications in the first 24 lm at 1200 °C up to 210 lm at 1650 °C.  Acknowledgements  TEM analyses were performed thanks to the financial support of grant N. FA8655-12-1-3004, with Dr. Ali Sayir as contract monitor. D. Dalle Fabbriche is greatly acknowledged for thermal treatments, C. Melandri for mechanical tests and S. Guicciardi for useful discussion on mechanical properties.  References  J. Routbort, F.  [1] E. Wuchina, M. Opeka, S. Causey, K. Buesking, J. Spain, A. Cull, Gutierrez-Mora, J. Mater. Sci. 39 (2004) 5939-5949. [2] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, J. Am. Ceram. Soc. 87 (2004) 1170-1172. [3] W.W. Wu, Y. Sakka, M. Estili, T.S. Suzuki, T. Nishimura, G.J. Zhang, Sci. Technol. Adv. Mater. 15 (2014) 014202. [4] D. Sciti, L. Silvestroni, C. Melandri, G. Celotti, S. Guicciardi, J. Am. Ceram. Soc. 91 (2008) 3285-3291. J. Zou, G.J. Zhang, C.F. Hu, T. Nishimura, Y. Sakka, J. Vleugels, O. Van der Biest, J. Am. Ceram. Soc. 95 (2012) 874-878. [6] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Am. Ceram. Soc. 91 (2008) 3530- 3535. [7] M. Kazemzadeh Dehdashti, W.G. Fahrenholtz, G.E. Hilmas, Corr. Sci. 80 (2014) 221-228. [8] D.G. Barton, S.L. Soled, G.D. Meitzner, G.A. Fuentes, E. (1999) 57-72.  [5]  Iglesia,  J. Catal. 181  \\x0c', 'L. Silvestroni, D. Sciti / Journal of Alloys and Compounds 602 (2014) 346-355  355  Intermetallics, VCH Publishers, Weinheim, New York, 1995. pp.  [9] G. Sauthoff, 115. [10] L. Silvestroni, D. Sciti, Effect of transition metal silicides on microstructure and mechanical properties of ultra-high temperature ceramics, in: J. Low, Y. Sakka, C. Hu (Eds.), MAX Phases and Ultra-High Temperature Ceramics for Extreme Environments, IGI Global, Hershey (PA), 2013, pp. 125-179. [11] D.G. Munz, J.L. Shannon Jr., R.T. Bubsey, Int. J. Fract. 16 (1980) R137-141. [12] L. Silvestroni, D. Sciti, J. Am. Ceram. Soc. 94 (2011) 1920-1930. [13] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Ceram. Soc. 89 (2006) 450-456. [14] J. Zou, G.J. Zhang, Y.M. Kan, J. Mater. Res. 24 (2009) 2428-2434. [15] H. Pastor, R. Meyer, Rev. Int. Haut. Temp. Refract. 2 (1974) 41-54.  J. Phase Equilib. Diff. 30 (2009) 564-  [16] Z. Guo, W. Yuan, Y. Sun, Z. Cai, Z. Qiao, 570. [17] E. Rudy, Part V. Phase Diagram W-B-C. AFML-TR-69-117 (1969). [18] L.Y. Chang, M.G. Scroger, B. Phillips, J. Am. Ceram. Soc. 50 (1967) 211-215. [19] L. Silvestroni, M. Nygren, D. Sciti, J. Eur. Ceram. Soc. 33 (2013) 2879-2888. [20] F.F. Lange, J. Am. Ceram. Soc. 63 (1980) 38-40. [21] A.G. Evans, A. Rana, Acta Metall. 28 (1979) 128-141. [22] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, J. Eur. Ceram. Soc. 27 (2007) 2495- 2501. [23] S.N. Karlsdottir, J.W. Halloran, J. Am. Ceram. Soc. 92 (2009) 481-486. [24] C.M. Carney, T.A. Parthasarathy, M.K. Cinibulk, J. Am. Ceram. Soc. 94 (2011) 2600-2607.  \\x0c']"
},{
  "_id": 258,
  "PDF": "TEM Study of the High‐Temperature Oxidation Behavior of Hot‐Pressed ZrB2–SiC Composites.pdf",
  "Text": "['TEM Study of  the High-Temperature Oxidation Behavior of Hot-Pressed ZrB2 -SiC Composites  Young-Hoon Seong, Seung Jun Lee, and Do Kyung Kim*,†  Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology (KAIST),  Yuseong-gu, Daejeon 305-701, Republic of Korea  The  oxidation behaviors of ZrB230 vol% SiC composites investigated at 1500°C in air and under were reducing condi−8 Pa, respections with oxygen partial pressures of 104 and 10 analyzed using  tively. The  oxidation  of ZrB2 transmission electron microscopy (TEM). Due to kinetic differ and SiC were  ence of oxidation behavior,  the three layers  (surface silica-rich  layer, oxide  layer, and unreacted layer) were observed over a  wide area of specimen in air, while the two layers (oxide layer,  and unreacted layer) were observed over a narrow area in spec imen under  reducing condition.  In oxide  layer,  the ZrB2 was into small grains  oxidized  to ZrO2 shape was also changed from faceted to round. This  accompanied  by  division  and the  layer also consisted of amorphous SiO2 with residual SiC and found dispersed in TEM. Based on TEM analysis of ZrB2-SiC composites tested under air and low oxygen partial pressure,  the ZrB2 begins to oxidize preferentially and the SiC remained without any changes at the interface between oxidized layer  and unreacted layer.  I.  Introduction  T HE  transition-metal  borides,  carbides,  and  nitrides  are  classified as ultra-high-temperature  ceramics  (UHTCs).  The UHTCs possess unique properties, (>3200°C),  including high melt ing  points  good mechanical properties, strong inertness.1-4 Because in UHTCs has increased in the field of aero oxidation  resistance,  and  chemical  of  this,  the interest  space applications,  for applications such as a thermal protec tion  systems  (TPS)  on  hypersonic  aerospace  vehicles  and  reusable  atmospheric reentry vehicles. Among the UHTCs, theoretical density (6.09 g/cm3), and has  ZrB2 has the lowest good thermal shock resistance because of (65-135 W/mK).5 These attributes could be for TPS and other aerospace applications.6-16 attempts have been made to enhance the oxidation  its high thermal  conductivity  advantageous  Many  resistance of ZrB2-based materials priate additives. The most common additive  through the use of appro is SiC, which  enhances the oxidation resistance via the formation of 1,16 as well as mechanical properties and sinterability.17-20 SiO2 The oxidation behaviors of ZrB2-SiC composites have been well-defined previously.21-24 When ZrB2-SiC composite is oxidizing environments (at a temperature  exposed 1500°C),  to  of  it  forms a layer structure consisting of  the following  (1) a continuous,  silica-rich layer,  (2) an SiC depletion layer,  and  (3)  an  unreacted  layer. This  surface  silica-rich  layer  prohibits  the transport of oxygen through the oxide scales, and allows the ZrB2-SiC composite to show parabolic mass gain kinetics.1,21  Under  reentry conditions, however, molecular oxygen will  be dissociated, due to the impact with the leading edge of the cap structure at hypersonic velocities.25,26  wing and/or nose  Therefore,  the O:O2 will be reduced during reentry. The oxidation mechanisms of  ratio (i.e.,  the oxygen partial pressure)  SiC in the UHTCs are different at different oxygen partial pressures.25 When the pO2 is high (>10 \\x005 Pa), SiC is oxidized and transformed to a viscous SiO2 phase at 1500°C. On the \\x005 Pa), SiC is low (<10 other hand, when the pO2 is transformed to a SiO phase with high vapor pressure 1500°C.25,27 Thus, along with temperature,  at  the oxygen partial  pressure  is an important parameter  in evaluating the oxida tion resistance of UHTCs. There are several methods to eval uate  the  high-temperature  oxidation  behaviors  of UHTCs.  Typically, high-temperature oxidation tests are conducted in  a furnace, which is a useful  tool  for high-temperature oxida tion behavior because of  the accurate control over the atmo sphere and temperature that  it allows.  There are many reports of  improved oxidation resistance,  which  have mainly  been  analyzed  using  SEM and  energy  dispersive X-ray  spectroscopy  (EDS). The  analysis  of  the  microstructural  shape  and phase  changes  in ZrB2 observation  and SiC  after  oxidation  is  also  important  and  of  the  ZrB2/ZrO2 and SiC/SiO2 our understanding of the oxidation behavior at high temper interfaces after oxidation improves  ature with varying pO2. There are few reports which observe the microstructure of oxidized ZrB2-SiC composites using transmission electron microscopy (TEM).28 L. F. He studied  that the kinetics of isothermal oxidation behavior Zr2Al3C4 at 500°C-1000°C and analyzed it using TEM.29 this article, we report an investigation of the oxidation of  of  In  SiC and ZrB2 and their oxidized phases after oxidation tests in air and at low oxygen partial pressure. In this research, we  aim to provide basic understanding of oxidation mechanism of ZrB2-SiC composite by TEM analysis.  II.  Experimental Procedures  (1)  Preparation  Commercially available raw powders were used in this study. a = b = 3.17 \\x17A, c = 3.53 A˚ ZrB2 powder (Hexagonal, , P6/mmm, size 3-5 lm, >99%, Grade A; H.C. Starck, Munich, a-SiC Germany) and powder (Hexagonal, a = b = 3.07 \\x17A, c = 15.08 \\x17A, P63mc, 98.5%, UF25; H.C. Starck) were used for hot-pressing. The  6H-polytype, 0.45 lm,  average  size  batches  consisted of 70 vol% of ZrB2 powder and 30 vol% of SiC powder. Many previous studies have reported that  ZrB230 vol% SiC composites showed improved sinterability and mechanical properties, and notably high-temperature resistance.8,22,30-34 70 vol% ZrB2 and 30 vol% SiC were used for tests in this study.  oxidation  Therefore,  composites  with  the oxidation  Before hot-pressing,  to reduce  the particle  size,  the ZrB2 raw powders were vibration milled for 30 min, using steel balls (~3 mm diameter spheres) and steel container; The median  particle  size  and  particle  size  distribution of  vibration-milled  W. Fahrenholtz—contributing editor  Manuscript No. 32016. Received September 8, 2012; approved February 1, 2013.  *Member, The American Ceramic Society.  †  Author to whom correspondence should be addressed. e-mail: dkkim@kaist.ac.kr  1570  J. Am. Ceram. Soc., 96 [5] 1570-1576 (2013)  DOI: 10.1111/jace.12246  © 2013 The American Ceramic Society  Journal  \\x0c', 'ZrB2 powder were D50 = 0.31 lm and D99/D50 = 1.99, while D50 = 4.16 lm D99/D50 = 2.76 and in as-received ZrB2 powder, respectively. During the vibration milling, 3.1 wt%  of Fe impurity was  introduced to ZrB2 powder due to wear of steel balls. Fe impurity was reduced to 0.042 wt% level by  acid treatment (3 M HCl, 1 h). The ZrB2 powder were subsequently mixed with 30 vol% SiC powder and milled for 24 h  in polyethylene bottles, with pure ethanol as (~4 mm  the solvent and  ZrO2 media. The powders were  balls  diameter  spheres)  as  the  milling  then carefully dried in a rotating  evaporator, to prevent (q = 6.09 g/cm3)  phase separation between the SiC (q = 3.21 g/cm3). After  ZrB2 drying,  and  the  the powders were  crushed with a mortar and sieved  using 325 mesh sieves.  The  powders were  densified  using  hot  pressing  (HP20 1000-3560; Thermal Technology Inc., CA), with a temperature of 1950°C and a pressure of 32 MPa applied for 2 h in a graphite furnace, under an argon atmosphere. The powders  were loaded into an 18-mm diameter boron nitride coated graphite die. The furnace was heated to 1800°C at a rate of 20°C/min, and then heated to 1950°C at a rate of 10°C/min with uniaxial pressure of 32 MPa. Above 800°C, both a ther mocouple  and a pyrometer  (Infrared Thermometer,  620 A;  Konica Minolta, Tokyo,  Japan) were used to monitor  the  temperature of the furnace and the graphite die. When the die temperature reached 1800°C, a uniaxial  load of 32 MPa  was applied. After 120 min,  the furnace was cooled to room 10°C/min, of and the load was temperature dropped below 1500°C. ~18 mm and a  temperature  at  a  rate  removed when the die  Specimens with a diameter of ~6.5 mm were fabricated, and samples with 4 mm 9 4 mm 9 3 mm were diced from the  thickness of  dimensions  of  specimens  for  the oxidation tests. The  samples prepared for  the oxidation  tests are listed in Table I.  (2)  Oxidation  The oxidation tests were performed using a horizontal  tube  furnace  equipped with a MoSi2 heating element. Before tests, specimens were prepared using conventional polishing with a diamond abrasive, down to a 1 lm finish. They were  the  then placed on an alumina boat filled with ZrO2 balls as a buffer layer, and inserted into the center of the furnace,  where  the  thermocouple was 1500°C for were conducted at 10 h, under \\x008 Pa) atmosphere an air (Z3S-H) and low oxygen partial pressure (pO2 = 10 (Z3S-L). The heating and cooling rates were both 5°C/min. To create low pO2 atmosphere conditions, CO gas containing 2000 ppm CO2 was used, which produced an oxygen partial \\x008 Pa at 1500°C. This oxygen partial pressure pressure of ~10 was selected based on the information obtained from a ther located. The  oxidation  tests  modynamic model, which predicted an oxygen partial pres sure in the SiC-depleted region during the oxidation of ZrB2-SiC composite system.25 The alumina (Al2O3, 99.9% purity, 60 mm outer diameter 9 5 mm wall thickness 9 1 m  the  length)  tube was used for CO/CO2 gas flowing and ends of  the tube were sealed using gastight end caps. A gas flow was maintained with a flow rate of ~100 cm3/min.  (3)  Characterization  The  densities  of  the  hot-pressed  specimens were measured  using the Archimedes method, and the  theoretical densities  of  the composites were calculated using the rule of mixture.  The microstructures of the cross-sections of both the as-made  and  oxidized  specimens were  characterized  using  scanning  electron microscopy (FE-SEM; Philips, XL30 FEG, Eindho ven, Netherlands). To analyze the microstructures of  the ver tical  sections before  and after  the  tests,  the  specimens  for  both cases were cross-sectioned and mounted in epoxy, carefully polished with a diamond abrasive down to a 1 lm fin ish,  and  cleaned  in  an  ultrasonic  bath with  acetone. The  thicknesses  of  the  resulting  reaction  layers were measured  from the polished cross-sections. The  tested specimens were  prepared  for  TEM  observations  using  ion  thinning  performed using  a  focused ion beam system (FIB; Quanta  3-D FEG, FEI, Eindhoven, Netherlands). Bright field (BF)  images,  SAED patterns,  high-resolution  TEM (HRTEM)  images, and EDS data were acquired using a transmission electron microscope (Tecnai G2 F30 S-twin, FEI, Eindhoven,  Netherland)  operating  at  300 kV. The  thermo-gravimetric  (TG) properties of  the ZrB2 and SiC starting powders were analyzed using a TG/DTA instrument (TGA 92-18; SETA RAM, Caluire, France). The 1400°C at a rate of 10°C/min.  temperature  was  raised  to  III.  Results and Discussion  (1)  Characterization of Sintered Composite  The bulk densities for the hot-pressed billets were measured, yielding densities of 5.22 and 5.25 g/cm3. The theoretical density of the ZrB230 vol% SiC composite was 5.23 g/cm3, determined using a rule of mixture calculation (6.09 g/cm3 for ZrB2, and 3.21 g/cm3 for SiC). Several assumptions for using a rule of mixture were used as follows: the 30 vol% SiC grains  were well dispersed in 70 vol% ZrB2 matrix and there are no reaction, substitution, and solid solution between ZrB2 and SiC. The relative densities of specimens were over 99.5% to  near  theoretical density. This  indicated that  the porosity did  not have a significant effect on the oxidation behavior.  Figure 1(a) shows a bright field TEM image of the hot-pressed ZrB2-SiC composite. The black grains in the figure are ZrB2, and the white grains are SiC. Intergranular faceted SiC grains with a grain size of 0.5-2.0 lm were observed at the  ZrB2  triple  and quadruple  junctions. The ZrB2  grains  also  Table I.  Summary of ZrB2-SiC Specimens: Compositions, Designations, Relative Density, and Conditions in the  Oxidation Tests  Composition (volume ratio)  ZrB2:SiC = 7:3  Designations  Z3S-H  Z3S-L  Relative  density (%)  99.5  99.8  Oxidation  test  conditions  Oxygen partial  pressure (Pa)  2 9 104 Air,  high pO2  \\x008 CO/ 10 CO2, low pO2  Temperature (°C)  1500  Time (h)  10  (a)  (b)  Fig. 1.  Typical  BF  image  of  as-sintered  ZrB2located at  30 vol% SiC  composite.  Intergranular SiC grains were  the  triple  and  quadruple  junctions of  the ZrB2 grains. The showing [01\\x161\\x161] taken from circled areas are ZrB2 and [0\\x16121] zone axis pattern from SiC.  inset SAED patterns  zone axis pattern from  May 2013  High-Temperature Oxidation of ZrB2-SiC  1571  \\x0c', 'showed a faceted shape, with grain sizes of 3-6 lm. The inset taken from the circled areas are [01\\x161\\x161] zone SAED patterns axis pattern of hexagonal ZrB2 and [0\\x16221] zone axis pattern of SiC-4H polytype (hexagonal structure) in Fig. 1(a). The  SiC-6H starting powder might partially transform to 4H-type  during sintering and the most stable polytypes carbide at 1800°C-2000°C were 4H and 6H.35  of  silicon  (2)  Oxidation Tests  Figure 2 shows the reacted depth as a function of for the ZrB2-SiC specimens oxidized in air and low \\x008 Pa) atmosphere at 1500°C. In air, the shape of parapO2 (10 bolic plots of oxidized depth versus time indicates that oxidation  exposure  time  follows  the  parabolic  rate  law. The  parabolic  rate law of of ZrB2-SiC It means the  oxidation  kinetics  implies  that  the  oxidation  composite  in  air  is  controlled  by  diffusion.  formed  silica-rich  oxide  scale  [reaction (1)]  is  protective  because  it  is dense and smooth as seen in Fig. 3(a). Under \\x008 Pa), low oxygen partial pressure (10 the oxidation kinetics at 1500°C is divided into two parts, parabolic kinetics at starting 3 h and linear kinetics thereafter. In other words,  the  oxidation kinetics deviates from the parabolic law and follow  a linear  law after 3 h of oxidation time. This  indicates  that  the rate-limiting of oxide growth changes  from the diffusion  of oxygen through the reaction product and (2)]; above 1300°C and 1100°C,  layer  [reactions (1)  respectively)  to reaction  between oxygen and SiC (reaction (3), SiO vaporizing below \\x0013 Pa at 1500°C), and SiO dissociation from pO2 ~8.8 9 10 \\x005 > pO2 > 8.8 9 10 \\x0013 Pa at SiO2 (reaction (4); range of 10 1500°C).25,36 parabolic-linear transition could be  The  explained as the result of the gradual interconnection of the such as pores crack and grain boundary.29 The conoxidized ZrB2-SiC grains loose and the number of grain-boundary interconnection is  defects,  nectivity  of  becomes  rough  and  increased. Due to low oxygen partial pressure,  the amount of (~38.1 Pa)  SiO, which  has  relatively  high  partial  pressure  from  SiO2 formation of pores. When the defects  increases  greatly,  possibly  resulting  in  the  interconnect with each  other  to form oxygen path (diffusion and transport)  chan nels,  the parabolic-linear transition of oxidation kinetics hap pens.  SiCðsÞ þ 3=2O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ  (1)  ZrB3 ðsÞ þ 5=2O2 ðgÞ ! ZrO2 ðsÞ þ B2O3  (2)  SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  (3)  SiO2 ðlÞ ! SiOðgÞ þ 1=2O2  (4)  Figure 3 shows cross-sectional SEM images of 1500°C  the materi als that were (2 9 104 Pa) and (b) under \\x008 Pa). The oxidation of 1500°C produced structures  oxidized  at  for  10 h  in  (a)  air  low oxygen  partial  pressure  (10  the Z3S-H [Fig. 3(a)] composites  at  consisting of four layers: layer, a ZrO2/ZrB2-SiC layer, and an unreacted layer. The thickness of layers (silica-rich layer + oxide layer) was 45 \\x06 5 lm. Several the oxidized previous studies reported that three layers (SiO2 layer (with SiO2 + ZrO2 cases), SiC-depleted layer, unreacted layer) were consistently formed after oxida1500°C.1,16,21,23-25 Because the ZrB2 phases oxidize rapidly, it results in the formation of a ZrO2 layer and silica-rich layer, via reactions (2) and (1), respectively.18  a surface  silica-rich layer, an oxide  layer:  observed  in  some  tion  tests  at  and  SiC  However, three different (SiO2 + ZrO2), this study; there was  layers  (silica-rich layer, oxide  layer  and  unreacted  layer) were  observed  in  no  SiC-depleted  layer  and  the  oxide  layer  consisted  of  ZrO2 (Therefore, we  and  amorphous  SiO2 layer  contained  unreacted  SiC.  call  this  “oxide  layer”  in this study) The formation of the SiCin the ZrB2-SiC system depends not only on the surrounding pressure and temperature conditions but also the SiC in the ZrB2 matrix.32 concretely, SiC-depleted layer is not formed  depleted layer  on the volume distribution of  In some  cases,  because of  the  volume  ratio of SiC/ZrB2, SiC distribution, partial pressure. The volume ratio of  and  internal  oxygen  SiC and ZrB2 the SiC fraction is  is 3:7 in this  study, and it  is  considered that  relatively high. The high fraction of SiC  increases  the degree of SiC interconnectivity. The pO2 value in this region cannot be defined clearly, but it is above \\x0013 Pa (boundary condition of reaction (3) 8.8 9 10 dent on the oxygen partial pressure), considering the pO2 for the ZrB2-ZrO2 equilibrium at 1500°C.36 Therefore, the SiC is oxidized to an amorphous SiO2 phase rapidly and it disperses 1500°C. The along the ZrO2 grain boundaries at reaction (1) was more dominant than reaction (3) in this region. Con is depen sequentially,  the SiO2 phase was  increased  by  the  internal  (b)  (a)  Fig. 2. Thickness of oxidation layer with time for exposure of ZrB230 vol% SiC at 1500°C (a) in air (2 9 104 Pa, Z3S-H) and (b) \\x008 Pa, Z3S-L). under low oxygen partial pressure (10  (a)  (b)  Fig. 3.  Cross-sectional SEM micrographs of oxidized ZrB230 vol% SiC composite (1500°C, 10 h). (a) Z3S-H and (b) Z3S-L showing a layered structure comprising an unreacted ZrB2-SiC layer (I), an oxide layer (II), and a uniform layer of SiO2 (III, only in Z3S-H). The oxidized layer in Z3S-H was thinner than that observed in Z3S-L.  1572  Journal of  the American Ceramic Society—Seong et al.  Vol. 96, No. 5  \\x0c', 'oxygen  partial  pressure  and  decreased  amount  of  SiO (g)  evaporation.  The SiO2 oxide scale on the surface is protective in an air atmosphere, because SiO2 is significantly less volatile than acts as a barrier  B2O3,  and  against  inward  oxygen  trans port.1,21,25  The volume increase upon ZrB2-SiC could also have been one of the amorphous SiO2 viscous flow to the silica-rich layer provides the passive oxidation behavior with  the  oxidation  of  the driving forces for surface.37 Thus,  a  a parabolic increase in oxidation depth. The oxide layer was  located underneath the  surface SiO2 region was similar to the original  layer,  and the micro structure of  this  structure,  because the tion.38 The  SiC was  removed  by  active  or  passive  oxida SiO2 might the ZrO2 boundaries under certain conditions (conditions for passive oxidation of SiC, temperature range from 1200°C to 1600°C pO2 > 8.8 9 10 between the unreacted layer and the oxide layer, even under \\x005 Pa), SiO high oxygen partial pressure conditions (pO2 > 10 (g) phase was transported to the surface via reaction (3),  have  remained  at  grain  and  above  \\x0013 Pa).25,32  At  the  interface  after  reaction (1)  occurred  near  the  interface  between  the  unreacted layer and the oxide  layer. This occurred because  this  region had  a  lower oxygen partial  pressure  compared  with that outside of  the specimen. However,  it consequently  showed  passive  oxidation  behavior  because  of  the  surface  SiO2 oxide scale that acted as a SiO (g) barrier layer. In contrast to Z3S-H, in Z3S-L, the silica-rich layer was  not  formed on the surface and, only two layers  (oxide layer,  unreacted layer) were observed. The depth of II of Fig. 3(b)] was 135 \\x06 3 lm, much thicker the oxidized than that measured for Z3S-H. A few studies have reported the ZrB2-SiC system under reducing conditions. Rezaie et al. explained that the high vapor  layers  [layer  on the oxidation behaviors  of  pressure of SiO (g) under reducing conditions leads to the active oxidation of SiC.25 The SiC is removed directly by active \\x0015 Pa < pO2 < 8.8 9 10 \\x0013 Pa) or it oxidizes to SiO2 [reaction (1)], which is then removed by \\x0013 Pa < pO2 < 10 \\x005 Pa) [reaction (4)]. volatilization (8.8 9 10 stable below pO2 ~ 1.9 9 10 \\x0011 Pa, or the ZrB2 will oxidize to form ZrO2 and B2O3. In other words, the oxidation behaviors of the ZrB2-SiC system depend on the precise pO2. The oxygen diffuses to the inside of the bulk more easily, due to the absence of a SiO2 layer on the surface and the smaller amounts of SiO2 at the grain boundaries under reducing conditions.  oxidation  (10  [reaction (3)],  Likewise,  is  (A) Oxidation 2 9 104 Pa):  Test  at High-pO2  Atmosphere  (Air,  Figure 4(a)  shows a typical BF image  taken  from the region slightly below the interface between the unre acted layer  (I) and the oxide  layer  (II)  in Fig. 3(a). Several  samples  were  prepared  to  find  this  interface  at  varying  depths.  In this  region,  some ZrB2 grains were oxidized and shapes were changed from faceted to uneven, while the  their  SiC grains  remained stable. Some parts of  the ZrB2 transformed. The ZrB2 therefore, divided into two grains: unreacted ZrB2 grain and oxidized ZrB2 (ZrO2 phase). It is well-known that below 1100°C, the oxidation of SiC (cr) is much slower than that of ZrB2,39 and that above 1100°C, SiC is oxidized rapidly via reaction (1) to form SiO2. At 1500°C, the oxidation the ZrB2 phase exhibited rapid linear kinetics. Even when the oxygen partial pressure of this region was much lower  grain  began to oxidize, and its phase was  grain was,  of  than that of  the surface,  the ZrB2 phase was oxidized preferentially, while the SiC phase was rarely oxidized. Figure 4(b)  shows  an HRTEM image  taken from the  squared area  in  Fig. 4(a),  at  the  interface  between  partially  oxidized ZrB2 arrangement of the left  and  unreacted ZrB2. The region clearly revealed it to be a hexagonal structure with a spacing of 3.5487 \\x17A for the {001} plane of The atomic arrangement  atomic  planar  the ZrB2 region also  phase.  of  the  right  clearly  showed structure  and planar  spacing. A monoclinic  ZrO2 right  structure was  revealed and the planar  spacing of  the  region did not match any of  the SiC polytypes. This  region had a planar spacing of 3.1639 \\x17A for the {\\x16111} plane of the monoclinic ZrO2 phase. There is a large difference in lattice volume at the interface between the ZrO2 (Vlattice = 140.7 A˚ in monoclinic ZrO2,) ZrB2 (Vlattice = 30.7 A˚ 3) grains. Therefore, stress might be concentrated at the interface between ZrO2 phase and ZrB2 phase. The deformation (transformation) also occurred to minimize  3  and  the  strain energy at  this  incoherent  interface. The  interface  between ZrO2 and ZrB2 might have moved leftward (in the direction of the inward ZrB2 grain) as the oxidation time increased. The partially oxidized ZrB2 firmed by element maps in Fig. 4(c). The Zr and O elements  grain was  also con were detected at  right  region which had a planar  spacing of  the monoclinic ZrO2 phase in ZrB2 grain and grain boundary while Si and C elements were only detected in SiC grain.  From the  results  of Fig. 4,  it  seems  the  oxygen might  be  diffused and oxidized along the ZrB2 grain boundary preferentially and preceded to inward ZrB2 grain. The BF image of Fig. 5 shows the microstructure of  the  oxide  interlayer  [layer  II  of Fig. 3(a)]  between  the  surface  SiO2 layer and the unreacted Z3S-H layer. It is closer to the outer silica-rich layer than to the un-oxidized ZrB2-SiC layer. The microstructure was different from that of the unreacted  layer, and the  shape of  the grains  changed from faceted to  round. Almost all of  the ZrB2 grains were oxidized to form an oxide phase (ZrO2); the grains were then divided into smaller grains with a size of 0.5-1.5 lm, and they changed  shape to minimize the surface energy. Many grain boundaries  were  created and the number of oxygen diffusion path was  increased due to the oxidation of ZrB2. In addition, the number of oxygen vacancies in the non-stoichiometric zirconium oxides (ZrO2\\x00d) might pressure.40 They accelerated the oxidation of  increase due to the low oxygen partial  the  specimen,  until  the  surface was  covered with  the  SiO2  amorphous  (a)  (c)  (b)  Fig. 4.  (a) BF image of  the  top part  in the unreacted ZrB2-SiC layer of Z3S-H, the ZrB2 grain and grain boundary begin to transform into a ZrO2 phase. (b) HRTEM image of the ZrB2/ZrO2 interface taken from the squared (c) The elements maps of partially oxidized ZrB2 grain  layer.  Immediately  below the  oxidized  area in (a).  and SiC grain.  May 2013  High-Temperature Oxidation of ZrB2-SiC  1573  \\x0c', 'phase. The  surface  of  SiC grain  also  oxidized  and  trans formed to amorphous phase of SiO2 with volumetric increase and viscous flow. Unreacted SiC remained at the midmost of  SiO2 and showed island structure. The SiC grains might also have been oxidized and trans formed  to  give  an  amorphous  SiO2 viscously into the grain boundaries. Based on the volatility diagram,25,38,41 the pO2 in this region was not allow the active oxidizing reaction [reaction (3)]. The pO2 for \\x0013 Pa, this region was above 8.8 9 10 considering the pO2 the ZrB2-ZrO2 equilibrium at 1500°C. \\x0013 Pa, the pO2 value the ZrB2-SiC lower than 8.8 9 10 the SiC in would oxidize to form SiO (g), and a layer of SiO2 (l) would not form in this region. A slight pO2 gradient was likely to exist across the layers.  phase with  flowing  low enough to  for  If  was  (B) \\x008 Pa): Oxidation  Test  at  Low-pO2  Atmosphere  (10  Figure 6 shows an STEM image of the specimen  tested at  low oxygen partial pressure;  the  image was  taken  from the  interface between the unreacted layer  (II) and the  oxidized layer  [layer I of Fig. 3(b)], which was thinned using  FIB  before  TEM analysis. Although  the  analyzed  area  was much narrower than that of Z3S-H, two small ZrO2/ZrB2+SiC/SiO2 (ZrO2/ZrB2-SiC section were clearly observed in one specimen (except the unreacted  layers  and  section)  ZrB2-SiC section), while similar a wide area in case of Z3S-H (Figs. 4 and 5). Because oxidation behavior of ZrB2-SiC at 1500°C under low pO2 kinetically much active. Section 1 of Fig. 6 shows the unre two sections were seen over  the  is  acted region with faceted ZrB2, and SiC grains. In section 2, some ZrB2 grains are oxidized and their shapes are changed, but the SiC grains are stable. Section 3 contains ZrB2 and SiC grains with ZrO2 grains and SiO2 amorphous phase. Figure 7(a) shows a high-magnification BF image taken  from section 2  (the  interlayer between the unreacted layer  and the oxide  layer)  in Fig. 6. ZrB2, ZrO2, and SiC grains Some ZrB2 grains were oxidized and their changed, and  were  observed.  shapes were  some ZrB2 small ZrO2 grains after the oxidation; in this region. Figure 7(b)  grains were  divided  into several  the SiC  was  stable  shows  an HRTEM  image taken from the circled area in Fig. 7(a). The interface  between  the  fully  oxidized ZrO2 observed. The atomic  grain  and  the  unreacted  ZrB2 grain on the left side of circled area clearly revealed a monoclinic structure with a planar spacing of 3.71 \\x17A for the {011} plane of the ZrO2 phase. The atomic arrangement of the grain on the right  grain was  arrangement  of  the  side of  circled area  also clearly  showed  planar spacing and its structure. A hexagonal structure and a  ZrB2 phase were revealed. This grain had a planar spacing of 3.58 \\x17A for the {001} plane of the ZrB2 phase. Figure 8 shows a high-magnification BF image taken from  section 3 of Fig. 6. The amorphous SiO2 phase was dispersed along the grain boundaries between the ZrO2 grains and the unreacted SiC grains, or at the SiC/SiC grain boundaries.  The SiO2 phase flowed viscously,  and wetted the SiC and  Fig. 5.  BF image of  the  interlayer between the  surface SiO2 glass phases (ZrO2 and amorphous SiO2) were mostly observed, and the unreacted ZrB2 and SiC grains remained in this layer.  layer  and  the  unreacted  layer.  The  oxide  Fig. 6.  STEM image of  the interface between the unreacted boride  layer  (Fig. 3(b)-I)  and the oxidized layer (Fig. 3 (b)-II) in Z3S-L. All of the sections (1: unreacted section, 2: ZrO2/ZrB2-SiC section, 3: ZrO2/ZrB2 + SiC/SiO2 section) existed in one specimen of Z3S-L.  and  (a)  (b)  Fig. 7.  (a) TEM BF image of  the ZrO2-SiC layer. Fig. 5) The ZrB2 grains and grain boundaries began to oxidize and transform into ZrO2, while the SiC grain maintained its phase. (b) HRTEM image and d-spacing  (Section 2 of  values  for  the ZrB2/ZrO2  interface  from the circled area in (a).  1574  Journal of  the American Ceramic Society—Seong et al.  Vol. 96, No. 5  \\x0c', 'ZrO2 formed to amorphous phase.  grains  after  the SiC grains were oxidized and trans In both cases  (air and reducing conditions),  the oxidation  started from the formation of ZrO2 at oxidized layer and unreacted layer. When ZrB2 grain boundaries of ZrB2 started to transform to ZrO2 phase first. Then the oxidation proceeded through grain boundaries  the interface between  is oxidized,  with oxygen diffusion and finally  the oxidation extened to  inward ZrB2 images in Figs. 5 and 7,  grains. This  phenomenon  can  be  seen  from  the ZrB2 grain is located inside and ZrO2 located outside of ZrB2 grain boundary. After that, the grains were divided into several small ZrO2 grains and round. As mentioned  is  the  shape  changed  from faceted  to  above,  the  oxygen might  transport  through  ZrB2 by rounded  grain  boundaries,  ZrO2 ZrO2 grains formation) and oxygen vacancies in the nonstoichiometric ZrO2\\x00d. Due to the multipath for diffusion of oxygen, the oxidation kinetics of ZrB2 was controlled by diffusion of oxygen through ZrB2 grain boundaries. The oxidation rate of SiC was relatively slower than that of ZrB2. The oxidation behavior of SiC started from the surface of  grain  boundaries  (created  SiC grain with transformation to amorphous SiO2 phase and the oxidation proceeded to inside the grain. When the SiO2 formed and covered on the surface of SiC grain, the oxygen  is difficult  to react with SiC because  surface SiO2 phase act as a barrier for oxygen diffusion. The surface SiO2 prohibited the oxygen diffusion inside and retarded the oxidation of  SiC. Therefore, unreacted SiC remained on the surface, SiO2 at the grain boundaries, and the oxidation kinetics of SiC  was  controlled by diffusion of oxygen at  surface SiO2,  i.e.,  diffusion-controlled kinetics.  Figure 9(a) shows a TEM BF image of  the oxide region in  the  oxidized  layer  [layer  II  of  Fig. 3(b)]  of  Z3S-L.  The  amorphous SiO2 phase was observed to disperse ZrO2 grain boundaries, including many pores. Figure 9(b) shows a magnified BF image of a grain boundary composed  along  the  of amorphous phase SiO2 and a pore in Fig. 9(a). The microstructure of this region was similar to that of the oxidized  layer  at  the  interface between the unreacted layer  and the  oxidized layer, as shown in Fig. 8 (section 3 of Fig. 6). How ever, many  pores were  observed  in  dispersed  SiO2 at the grain boundaries with absence of SiC phase after the test— sufficiently oxidized—  despite  the  fact  that  the  region was  indicated that  the oxygen partial pressure of  this  region was  not  low enough to rapidly evaporate the SiO phase from the  SiO2 phase (pSiO was low at this pO2). Based on the thermodynamic calculations and the volatility diagram, the range of  oxygen partial pressure  that allowed active oxidation in the \\x0015 Pa \\x005 Pa.42 The from 10 to 10 from the SiC phase in the  volatile SiO (g) phase  is  SiO phase  evaporated  directly  range of oxygen partial pressure \\x0013 Pa reaction (3), whereas 8.8 9 10 amorphous phase reaction (1)  from  \\x0015 Pa 10 the SiC transformed to  to  a  SiO2 phase evaporated from the SiO2 phase with a relatively low partial vapor pressure after oxidation in the range of oxygen \\x0013 Pa \\x005 Pa 8.8 9 10 partial pressure from to 10 reaction (4).25 The amorphous SiO2 phase therefore remained at the grain boundaries with pores, because the oxygen partial \\x008 Pa in this for Z3S-L was 10 the amorphous SiO2 phase originated from the the SiO (g) phase, and using pO2 = 10 \\x008 Pa, of can be used to calculate pSiO = 38.15 Pa. This showed that the ZrB2-SiC composite under the oxidation kinetics of partial pressure of oxygen at 1500°C is different from those  and  then,  the  SiO  pressure  study. The pores  in  evaporation  reaction (4)  low  of  the  composite  in ambient pressures, and the  existence of  protective layer on the surface has an important ZrB2-SiC shows a highly magnified BF image of amorphous SiO2 at the grain boundary taken from the squared area in  role to the  oxidation  kinetics  of  composites.  Figure 9(c)  Fig. 9(b). There were many ~5 nm. Also,  particles  in  amorphous  SiO2 atomic arrangement was 2.071 A˚ {301}  with  size  of  the  observed with  a  planar  spacing  of  for  the  plane of  the ZrSiO4 phase. in Fig. 9(d) and the Zr, Si,  It  is confirmed by element maps  and O were detected as parti cles. This  result was  caused  by  further  reaction  between  ZrO2 and dissolves  and  amorphous  SiO2. The into crystalline ZrO2 until reached when ZrO2 and amorphous silica coexisted, ter by precipitation of ZrSiO4.43  interstitial  silicon  diffuses  the solution limit  is  thereaf IV.  Conclusions  The ZrB2-30 vol% SiC in air (pO2 = 104 Pa), (pO2 = and under reducing conditions \\x008 Pa) at 1500°C for 10 h. The microstructures and oxidation 10 depths of the specimens were observed using SEM and phase  composites  were  oxidized  transformation and microstructure of  the grains/grain bound aries on each layer were analyzed using TEM.  Fig. 8.  TEM BF image of  the  reacted oxide  layer.  (Section 3 of  Fig. 5) The ZrO2 and SiC grains were observed and amorphous SiO2 was dispersed at grain boundary.  (a)  (b)  (d)  (c)  Fig. 9.  (a) TEM BF image of  the reacted oxide layer  (II)  in Z3SL;  the amorphous SiO2 dispersed along the ZrO2 grain boundaries with pores. (b) The magnified image of (a). The nano-sized particles  were  observed  (black  spots)  in  amorphous  SiO2. particles  (c) The  highly  magnified  image  of  particles  in  (b).  These  have  been  identified as ZrSiO4 by atomic arrangement with a planar spacing of 2.071 \\x17A for the {301} plane of the ZrSiO4 phase and (d) Zr, Si, O element maps.  May 2013  High-Temperature Oxidation of ZrB2-SiC  1575  \\x0c', 'Based on TEM results,  the three layers (surface SiO2 layer, layer, and unreacted layers) were observed in Z3S-H  oxide  and the  two layers  (oxide  layer,  and unreacted layer) were  observed in Z3S-L with varying depths after oxidation test.  The  SiO2 in Z3S-H, because of  and  residual  SiC were  dispersed  in whole  oxide  layer  the structural distribution of SiC  in the ZrB2 matrix and internal oxygen partial pressure. contrast, active oxidation behavior and no surface SiO2 layer were observed in Z3S-L, and the amorphous SiO2 phase also remained at the ZrO2 grain boundaries in Z3S-L. The results from the TEM analysis, in both  In  cases  (air  and  reducing  conditions),  ZrB2 was firstly and then,  oxidized  and  trans formed  to ZrO2 interface between unreacted layer and oxidized layer.  phase  SiC was  oxidized  at  the  The  oxidation  and  transformation  of  ZrB2 and the oxidation proceeded to the  was  started  from grain boundaries  inside  of  grain  which  showed  outside  ZrO2 divided  and  inside  ZrB2 ZrO2 from facet  structure. Then,  the  grains were  into  several  grains  after  fully  oxidizing with  the  shape  changing  to round. The SiC started to oxidize and trans form into SiO2 the SiO2 was dispersed in grain boundaries layer of composite due to high viscosity  from the  surface of SiC grain. After  that,  in whole oxide  and  volumetric  increase.  The  unreacted  SiC  existed  in  amorphous  SiO2  which has an island structure.  The oxidation kinetics of ZrB2 might be controlled by O2 diffusion and transport through the ZrB2 grain boundaries and ZrO2 grain boundaries, respectively, and the oxidation kinetics of SiC might be controlled by O2 diffusion through SiO2 because surface SiO2 acted as an oxygen diffusion barrier. The behavior  oxidation  in  structural  changes was  similar, but  the oxidation kinetics was different. TEM analysis  is one of the good approaches for behaviors of ZrB2-SiC-based UHTCs.  understanding oxidation  Acknowledgments  This work was  supported  by Defense Acquisition Program Administration  and Agency for Defense Development under  the  contract UD110093CD and  the Priority Research Centers Program through the NRF funded by MEST  (2009-0094041).  References  1W. G. Fahrenholtz, G. E. Hilmas,  I. G. Talmy,  and  J. A. Zaykoski,  “Refractory Diborides of Zirconium and Hafnium,” J. Am. Ceram. Soc., 90 [5] 1347-64 (2007). 2N. P. 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  "_id": 259,
  "PDF": "Textured HfB2-based ultrahigh-temperature ceramics with anisotropic oxidation behavior.pdf",
  "Text": "['Available online at www.sciencedirect.com  Scripta Materialia 60 (2009) 913-916  www.elsevier.com/locate/scriptamat  Textured HfB2-based ultrahigh-temperature ceramics with anisotropic oxidation behavior  De-Wei Ni,a,b Guo-Jun Zhang,a,* Yan-Mei Kana and Yoshio Sakkac,*  aState Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics,  bGraduate School of the Chinese Academy of Sciences, Beijing 200049, China cNano Ceramics Center and World Premier International Research (WPI) Center Initiative for  Shanghai 200050, China  Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), Tsukuba, Ibaraki 305-0047, Japan  Received 15 January 2009; revised 4 February 2009; accepted 7 February 2009  Available online 13 February 2009  Based on recent preliminary results of textured ZrB2-based ceramics, c-axis oriented HfB2-based ultrahigh-temperature ceramics with a Lotgering orientation factor as high as 0.91 were prepared by slip casting in a strong magnetic field alignment, followed by spark plasma sintering. The textured sample displays significantly anisotropic properties. Compared with textured ZrB2-based ceramics, the HfB2-based material shows much better oxidation resistance.  Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: Hafnium diboride; Slip casting; Strong magnetic field alignment (SMFA); Texture; Anisotropy  Transition metal borides, mainly hafnium diboride (HfB2) and zirconium diboride (ZrB2), belonging to the class of ultrahigh-temperature ceramics (UHTCs), are currently expected to be a potential candidate material for aerospace applications because of their high melting points (>3000 °C) and excellent chemical stability. The introduction of SiC improves the properties, especially the oxidation resistance, of these diborides [1,2]. Despite their potential, their widescale use still presents a challenge, due, in part, to limitations in their high-temperature properties and oxidation resistance. To improve their performance, most research has focused on composition design of existing materials - i.e. selecting different additives to change the material composition to improve the performance of materials [3-5]. However, grain boundary phases deriving from the additions often drastically deteriorate the strength and other mechanical properties at elevated temperature [6,7]. The controlled development of microstructures has recently become an important topic in ceramic processing because it allows improvement in mechanical, thermal, electrical and other properties. As a common  * Corresponding  authors; e-mail  addresses:  gjzhang@mail.sic.ac.cn;  sakka.yoshio@nims.go.jp  method for microstructure tailoring, texture development offers the unique opportunity to optimize the performance of ceramics, and has been widely used in the materials field [8-10]. ‘‘Templated grain growth” (TGG) and ‘‘hot-working” methods [8,11] are two major methods for texture development. However, to the knowledge of the authors, virtually nothing has been published in the open literature discussing the texture development of UHTCs, partly due to the difficulty of preparing templated grains. Recently, strong magnetic field alignment (SMFA) has been successfully utilized to texture ‘‘non-magnetic” ceramics (e.g. Al2O3, TiO2, AlN and Si3N4); this technique involves the alignment typically P10 T, of particles in a strong magnetic field, during slurry consolidation, followed by sintering [12-15]. The alignment is achieved with the axis showing the highest magnetic susceptibility parallel to the magnetic field. In this novel method, the primary requirement is that the target material should be non-cubic (or anisotropically magnetic) and the suspension should be well dispersed with low viscosity. Compared with the conventional methods for texture development, SMFA has the attraction that it is not limited by particle morphology. Using the SMFA method, we recently successfully prepared c-axis highly textured ZrB2-based ceramics,  1359-6462/$ see front matter Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2009.02.013  \\x0c', '914  D.-W. Ni et al. / Scripta Materialia 60 (2009) 913-916  which showed anisotropic properties [16]. As another important candidate material for aerospace applications, HfB2-based composites appear more promising compared to ZrB2-based composites; the Hf materials have higher melting point, hardness and thermal conductivity, etc. [17]. HfB2 has the same hexagonal structure as ZrB2; accordingly, the SMFA method should be effective for texturing HfB2-based ceramics as well. This paper describes the texture development of HfB2based ceramics via SMFA and discussion is focused on the anisotropic oxidation behavior. Commercially available SiC (D50 = 0.45 lm, purity: 98%) and synthesized HfB2 were used as starting materials; HfB2 powder was synthesized via modified carbothermal/borothermal reduction of HfO2 at a relatively low temperature (1600 °C) [18]. In this study, HfB2, HfB2-5 vol.% SiC and HfB2-20 vol.% SiC composites were prepared, and are designated as HB, HS5 and HS20, respectively. Distilled water was used as a dispersing medium and 1.5 wt.% (based on the total weight of the powders) polyethylenimine (PEI, M w = 10,000, Wako Pure Chemical Industries Ltd., Tokyo, Japan) was chosen as a dispersant. First, the synthesized HfB2 powder was planetary milled for 8 h in a Si3N4 jar with Si3N4 balls to break up the agglomerates. Slurries with solid loading of 30 vol.% were then prepared by ball milling for 24 h in a plastic bottle with Si3N4 balls. After degassing, slip casting was performed using a glass case set on a plaster block with a 0.2 lm membrane filter in a 12 T magnetic field. The direction of the magnetic field was perpendicular to the direction of slip casting. After consolidation, the drying was performed in an oven at 50 °C for 24 h, followed by cold isostatic pressing at 392 MPa. The green bodies were calcined at 700 °C for 0.5 h in vacuum to burn out the polymer dispersant before sintering. The sintering was conducted in a spark plasma sintering (SPS) furnace at 1900 °C under a pressure of 50 MPa in vac uum. The final density was measured using the Archimedes method. The crystallographic orientation and phase compositions were evaluated by X-ray diffraction (XRD, JDX 3500, Rigaku Co., Japan) using CuKa radiation (k = 1.54178 A˚ ) and a scanning rate of 2 min\\x001 on different surfaces of the samples. The Lotgering orientation factor, f(0 0 l), was used to evaluate the degree of texture in the green and sintered samples [16]. The hardness and fracture toughness were measured by the indentation method [19] using a load of 5 kg for 10 s on a polished surface and the reported value was an average of three measurements. After grinding to a 3 lm surface finish, the oxidation behavior of HS20 was studied in a box furnace. The sample was placed on a zirconia plate with minimal contact area and the two surfaces of SS and TS (defined below) were kept under the same environmental conditions. The furnace was heated at \\x1810 °C min\\x001 to 1600 °C and held for 10 min- 10 h in stagnant air. The microstructures were observed by scanning electron microscopy (SEM, JSM-6500, Hitachi Co., Tokyo, Japan) along with energy-dispersive spectroscopy (EDS) for chemical analysis. After planetary milling, the HfB2 powder was welldispersed and had an average particle size of 0.67 lm. The rheological behavior analysis showed that the 30 vol.% HfB2 aqueous slurries almost maintained a constant viscosity of \\x1820 mPa s as the shear rate increased from 25 to 400 rpm and no shear thinning behavior appeared. Thus, the results confirmed that 1.5 wt.% PEI could stabilize the HfB2 aqueous suspension efficiently. The green bodies prepared by slip casting in the magnetic field showed obvious c-axis orientation, as revealed by the XRD patterns collected from the top and side surfaces of the green bodies (Fig. 1). The top and side surfaces, which are parallel and perpendicular to the magnetic direction, are denoted as TS  Figure 1. XRD patterns collected on the TS and SS surfaces of the green bodies: (a) HB;  (b) HS5; (c) HS20.  \\x0c', 'D.-W. Ni et al. / Scripta Materialia 60 (2009) 913-916  915  and SS, respectively. It is clear that the diffraction peaks of all (0 0 l) planes, typical of (0 0 1) and (0 0 2) planes perpendicular to the c-axis of the unit cell, show much stronger relative intensities on the SS surface than on the TS surface. However, the diffraction peaks of the (h k 0) planes parallel to the c-axis show relatively weak intensities on the SS surface, but become stronger on the TS surface. This indicates that the presence of a magnetic field leads to the strong crystallographic orientation shown by the alignment of the c-axis of the HfB2 crystals parallel to the magnetic field due to the anisotropic magnetic susceptibility of vc > va,b, which is very similar to the result of ZrB2 in our previous work [16]. Based on the XRD peak intensity data on the SS surface, the Lotgering orientation factor f(0 0 l) was determined to be 0.61, 0.62 and 0.69, respectively, for the three samples, suggesting a strong texture of HfB2 grains, while some HfB2 particles were still not oriented. One reason for this is probably associated with the existence of agglomeration in HfB2 powder. On the other hand, as the HfB2 powder was synthesized at a relatively (1600 °C) low temperature [18], crystallization of HfB2 might be incomplete and have resulted in a partially orientated state. With the addition of SiC, the orientation of HfB2 grains increased slightly; the reason for this is not clear but might due to improved breaking of the agglomerates, resulting in a more uniform distribution of the slurries. The typical physical and mechanical properties of the SPS samples are presented in Table 1. It is clear that the relative density of all the samples, except HB, is higher than 96% after SPS at 1900 °C/2 min/50 MPa. For comparison, HB was also prepared via the same process in the absence of a magnetic field, and a theoretical density of \\x1895% was obtained. This result is very similar to the textured ZrB2 reported in our previous work [16] and indicates that it is very difficult to densify orientated monolithic MB2, which may be due to the thermal expansion and elastic constants anisotropy of MB2 grains. As can be seen from the results listed in Table 1, the textured samples show anisotropic mechanical properties. The hardness on the SS surface is about 10% higher than that on the TS surface for both HS5 and HS20, due to the intrinsic anisotropic properties of HfB2 [20] and the residual stress resulting from the coefficient of thermal expansion anisotropy along the aand c-axis directions of HfB2. However, fracture toughness anisotropy in different directions is less evident. During sintering, some poorly crystallized HfB2 particles not oriented during slip casting could crystallize completely or grow into the oriented HfB2 grains by dif Table 1. Mechanical properties of the SPS samples.  Properties True density (g cm\\x003)  Relative density (%)  Apparent porosity (%)  Vickers’ hardness  (Hv5) (GPa)  Fracture toughness KIc (MPa m1/2)  HB  9.53  85  1.15  HS5  10.55  97.6  0.07  HS20  9.45  98.3  0.01  SS  TS  SS  TS  -  -  -  -  20.6 ± 0.6  18.9 ± 0.4  5.8 ± 0.3  5.5 ± 0.5  21.1 ± 0.5  19.3 ± 0.8  5.9 ± 0.2  5.5 ± 0.4  Figure 2. SEM images of HS5 on different polished surfaces.  fusion, resulting in a Lotgering orientation factor f(0 0 l) as high as 0.91 for the SPS samples. Despite the high degree of orientation, we could not find any obvious morphology difference between the SS and TS surfaces (Fig. 2), owing to the HfB2 grains not being elongated. If the oriented grains were elongated, they could act as reinforcements and increase the opportunity for crack bridging and deflection, thereby resulting in both high strength and high fracture toughness in the direction perpendicular to the grain alignment as demonstrated in Si3N4 ceramics [9]. This should be an important subject for future work. It has been widely recognized that oxidation of MB2- SiC (M = Zr and Hf) at \\x181500 °C usually produces a layered structure [21-23]. No obvious oxidation is found on both surfaces (SS and TS) of HS20 after holding at 1600 °C for 10 min in stagnant air, which hence shows much better oxidation resistance than textured ZrB2- 20 vol.% SiC [16]. As shown in Figure 3, oxidation of HS20 at 1600 °C for 10 h produces a layer structure with a homogeneous thickness distribution on each oxidized surface. However, the oxidation layer formed on the SS surface differs significantly from that on the TS surface, indicating anisotropic oxidation behavior of the textured samples. On the SS surface, the oxidation layer consists of: (1) a SiO2-rich glassy layer (\\x1860 lm); and (2) a SiCdepleted layer composed of HfB2 (>80 lm) (due to limitations of the magnification used, we could not see the edge of the SiC-depleted layer). While on the TS surface, the oxidation layer consists of: (1) a SiO2-rich glassy layer (<20 lm); (2) a transition layer composed of HfO2 and SiO2 (\\x1860 lm); and (3) a SiC-depleted layer composed of HfB2 (\\x1830 lm). One reason for the anisotropic oxidation behavior might be the intrinsic characteristic of crystalline HfB2. On the other hand, the mass transmission (such as oxygen, SiO, etc.) might be also anisotropic, resulting in the SiO2-rich glassy layer formed on SS surface being markedly thicker, thereby providing a much better protective effect on the matrix. Conversely, the thin SiO2-rich glassy layer on the TS surface might result in a relatively high oxygen content in the next layer beneath it and consequently caused the oxidation of HfB2. The details on the oxidation anisotropy are under investigation. Moreover, there is a white thin layer in the SiO2-rich glassy layer indicated by arrows in Figure 3a and e. It is difficult to identify its composition due to the low sensitivity of energy-dispersive spectroscopy. Wavelength dispersive X-ray spectroscopy (JEOL JXA-8100F, Japan) analysis indicates that the white thin layer in SiO2-rich glassy layer is composed of HfO2, probably the result of oxidation of HfB2 at the very initial stage during heating.  \\x0c', '916  D.-W. Ni et al. / Scripta Materialia 60 (2009) 913-916  Financial support from the Chinese Academy of Sciences under the Program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program), the National Natural Science Foundation of China (No. 50632070) and the Science and Technology Commission of Shanghai (No. 08520707800) are greatly appreciated. This study was supported in partly by the Grant-in-Aid for Scientific Research of the JSPS and World Premier International Research (WPI) Center Initiative on Materials Nanoarchitronics (MANA), MEXT, Japan. The authors are grateful to Drs. T.S. Suzuki, T. Uchikoshi, S. Gustavo and G. Salvatore for help in experiments and useful discussions.  P.L. Wang,  Scripta  J. Am. Ceram. Soc. 81  [1] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, J. Eur. Ceram. Soc. 19 (1999) 2405. [2] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, J. Am. Ceram. Soc. 87 (2004) 1170. [3] L. Silvestroni, D. Sciti, Scripta Mater. 57 (2007) 165. [4] W.W. Wu, G.J. Zhang, Y.M. Kan, P.L. Wang, K. Vanmeensel, J. Vleugels, O. Van der Biest, Scripta Mater. 57 (2007) 317. [5] J. Zou, G.J. Zhang, Y.M. Kan, Mater. 59 (2008) 309. [6] F. Monteverde, S. Guicciardi, A. Bellosi, Mater. Sci. Eng. A 346 (2003) 310. [7] F. Monteverde, A. Bellosi, J. Mater. Res. 19 (2004) 3576. [8] K. Hirao, M. Ohashi, M.E. Brito, S. Kanzaki, J. Am. Ceram. Soc. 78 (1995) 1687. [9] N. Kondo, T. Ohji, F. Wakai, (1998) 713. [10] H. Imamura, K. Hirao, M.E. Brito, M. Toriyama, S. Kanzaki, J. Am. Ceram. Soc. 83 (2000) 495. [11] N. Kondo, Y. Suzuki, T. Ohji, J. Am. Ceram. Soc. 82 (1999) 1067. [12] T.S. Suzuki, Y. Sakka, K. Kitazawa, Adv. Eng. Mater. 3 (2001) 490. [13] Y. Sakka, T.S. Suzuki, J. Ceram. Soc. Jpn. 113 (2005) 26. [14] X.W. Zhu, T.S. Suzuki, T. Uchikoshi, T. Nishimura, Y. Sakka, J. Ceram. Soc. Jpn. 114 (2006) 979. [15] X.W. Zhu, T.S. Suzuki, T. Uchikoshi, T. Nishimura, Y. Sakka, J. Am.Ceram. Soc. 91 (2008) 620. [16] D.W. Ni, G.J. Zhang, Y.M. Kan, Y. Sakka, Scripta Mater. (2009), doi:10.1016/j.scriptamat.2008.12.027. [17] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, J. Am. Ceram. Soc. 90 (2007) 1347. [18] D.W. Ni, G.J. Zhang, Y.M. Kan, P.L. Wang, Ceram. Soc. 91 (2008) 2709. [19] A.G. Evans, E.A. Charles, J. Am. Ceram. Soc. 59 (1976) 371. [20] S. Otania, M.M. Korsukova, T. Aizawa, J. Alloys Compd. (2008), doi:10.1016/j.jallcom.2008.10.094. [21] A. Chamberlain, W. Fahrenholtz, G. Hilmas, D. Ellerby, Refract. Appl. Trans. 1 (2005) 1. [22] W.G. Fahrenholtz, J. Am. Ceram. Soc. 90 (2007) 143. [23] W.M. Guo, X.J. Zhou, G.J. Zhang, Y.M. Kan, Y.G. Li, P.L. Wang, J. Alloys Compd. (2008), doi:10.1016/ j.jallcom.2008.02.108.  J. Am.  Figure 3. Cross-sectional images (a) and (e) and compositional maps the polished surfaces of HS20 after oxidation at 1600 °C for 10 h:  of  (a)-(d) oxidation results for the SS surface; (e)-(g) oxidation results for  the TS surface.  In summary, highly c-axis oriented HfB2-based ceramics with a Lotgering orientation factor f(0 0 l) of up to 0.91 were obtained by slip casting in a strong magnetic field alignment, followed by SPS. Due to the anisotropy of the thermal expansion and elastic constants of HfB2, it was very difficult to densify orientated monolithic HfB2. For the densified samples, it was found that hardness on the SS surface was superior to that on the TS surface. The oxidation resistance of the textured sample (HS20) also showed significant anisotropy. The SiO2-rich glassy layer formed on SS surface was markedly thicker, providing a much better protective effect on the matrix. It is therefore expected that the anisotropic properties of the textured materials will provide more opportunities for the material design and selection of UHTCs.  \\x0c']"
},{
  "_id": 260,
  "PDF": "The addition of lanthanum hexaboride to zirconium diboride for improved oxidation resistance.pdf",
  "Text": "['Scripta Materialia 57 (2007) 1036-1039  www.elsevier.com/locate/scriptamat  The addition of  lanthanum hexaboride to zirconium  diboride for improved oxidation resistance  Xing-hong Zhang, Ping Hu,* Jie-cai Han, Lin Xu and Song-he Meng  Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, China  Received 24 May 2007; revised 20 July 2007; accepted 22 July 2007  Available online 6 September 2007  The oxidation resistance of ZrB2-SiC-based ultrahigh-temperature ceramic (UHTC) is remarkably enhanced by the addition of lanthanum hexaboride, which has demonstrated outstanding oxidation resistance to temperatures up to 2400 °C. The improved oxidation resistance is attributed to the formation of a coherent compact scale, which acts as an effective barrier against the inward diffusion of oxygen.  Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Keywords: Oxidation resistance; Lanthanum hexaboride (LaB6); Zirconium diboride (ZrB2); Ultrahigh-temperature ceramic (UHTC)  Ultrahigh-temperature ceramics (UHTC), including some refractory metal diborides, have been studied and developed since the 1960s [1,2]. In recent years, great attention has been paid to ZrB2 and HfB2. Their striking thermochemical stability in terms of extremely high melting point, resistance to ablation/oxidation at high temperatures and thermal-shock resistance represent key requirements applicable to thermal protection systems for re-entry space vehicles with sharp leading-edge profiles [3-6]. ZrB2-SiC and HfB2-SiC are currently considered as the baseline UHTC. Indeed, varying the initial composition of these UHTC by changing the SiC content has given added flexibility in optimizing specific microstructural designs: adjusting the SiC content in ZrB2 and HfB2 matrices has proved beneficial for improving oxidation and ablation resistance, without being detrimental to high-temperature stability [4-6]. The significant improvement in oxidation resistance of ZrB2 (HfB2)-based UHTC below 1700 °C have been achieved by incorporation of silicides due to the formation of silica glass layer with low oxygen permeability, which provide an efficacious protective oxidation barrier. However, silicides become invalid at higher temperatures, especially above 2000 °C, because of active oxidation and evaporation. Thermodynamicsbased calculations have also shown the temperature limits in oxidizing environments for state-of-the-art,  * Corresponding  author.  Tel.:  +86  13845077148;  fax:  +86  45186402382; e-mail: huping@hit.edu.cn  high-temperature SiO2-forming materials, such as SiC, Si3N4 and MoSi2 [7]. Although great efforts have been made in this aspect, no effective approach has been reported up to now. This work highlights the significant improvement in the oxidation resistance of ZrB2-SiCbased UHTC with the addition of 10 vol.% LaB6. The role of LaB6 is also examined and discussed. The ZrB2 and LaB6 powders (Northwest Institute for Non-ferrous Metal Research, China) were of the same purity (99.5%) and both had a mean particle size of 5 lm. The SiC powder (Weifang Kaihua Micro-powder Co., Ltd., China) was predominantly a-SiC, and it had a reported purity of 99.5% and a mean particle size of 2 lm. Batches containing ZrB2-20 vol.% SiC-10 vol.% LaB6 were prepared. Powders were milled in a polymer-coated bucket charged with ethanol using WC balls for 8 h and dried in a rotating evaporator. Milled powders were uniaxially hot pressed in a boron nitride coated graphite die at 2000 °C for 60 min under vacuum (0.5 mbar) and 30 MPa of applied pressure. Specimens were first polished and ultrasonically cleaned by acetone and absolute ethanol. The final densities of composites were measured by Archimedes’ method, and the relative density was estimated by the rule of mixture. Three sample coupons of U18.8 · 14.2 mm were cut from hot-pressed plates for oxidation tests. The specimens were ultrasonically cleaned in acetone followed by absolute ethanol, dried and then weighed (to an accuracy of 0.1 mg). Oxidation tests were carried out with an oxyacetylene torch. Models were placed in graphite holders which enabled test durations in excess of 600 s.  1359-6462/$ see front matter Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  doi:10.1016/j.scriptamat.2007.07.036  \\x0c', 'X.-h. Zhang et al.  / Scripta Materialia 57 (2007) 1036-1039  1037  Oxidation tests were conducted according to GJB323A96 standard on the oxyacetylene ablation equipment [8]. The pressure and flux of acetylene were 0.1 MPa and 1.15 m3 h\\x001, and for oxygen 0.5 MPa and 2.30 m3 h\\x001, respectively. The emissivity was determined by a twocolor radiation pyrometer and the temperature of specimen surface was measured using a one-color pyrometer (Mikron Instrument Co., Inc., Oakland, NJ). Two ZrB2-20 vol.% SiC sample coupons were also tested under the same conditions for comparison. An influence of oxidation was investigated by measuring the mass change and thickness of the scale. Xray diffraction and scanning electron microscopy/energy dispersive spectroscopy (FEI Sirion, Holland) were used to characterize the phase composition and microstructure of the surface and cross section of the samples after oxidation experiments. The bulk density of the hot-pressed material was 5.93 g cm\\x003, which corresponds to 99.2% of the theoretical density estimated with the rule of mixture. Figure 1 shows scanning electron micrographs of the polished surfaces of the ZrB2-SiC-LaB6. The microstructure of the composites is regular, with a mean grain size is about 5 lm, and the residual porosity is very scarce. The SiC particulates (dark phases) are distributed homogeneously in the diboride matrix and no agglomeration was detected. Figure 2 shows the photographs of ZrB2-SiC-LaB6 during (left) and after (right) oxyacetylene torch testing. The macromorphology of the oxidized surface of the specimen oxidized for 600 s with an oxyacetylene torch shows a mechanical scour in the surface center zone, resulting from blowing by the airflow. The surface layer appears compact and adherent, and no cracks or spallation were detected there. The surface temperature increased drastically in a very short time when the  samples were exposed to the oxyacetylene torch heater. Specifically, the surface temperature rose dramatically to 2000 °C and then began to rise gradually, reaching a maximum temperature of over 2400 °C, as can be seen in Figure 3. A high thermal stress will occur in the tested sample due to the large temperature gradient. However, no cracks were observed on the sample. This fact suggests that the composite exhibits good thermal-shock resistance. A total weight loss only of 0.2% was measured. The measured weight loss is primarily ascribed to the active oxidation of SiC and volatilization of some oxidation products such as B2O3 and SiO2. Figure 4 shows the macrograph and cross-sectional micrograph of oxidized ZrB2-20 vol.% SiC after testing under the same conditions. However, extensive damage was observed at the external surface and the outer oxide scale was abscised. Cross-sectional examination showed that the oxide scale is incompact and not adherent to unaltered material, as shown in Figure 4(b). Moreover, the total weight change for ZrB2-20SiC (3.1%) is far higher than that for ZrB2-20SiC-10LaB6 (0.2%). There are three kinds of oxidants, namely, H2O, CO2 and O2, in the present case. H2O and CO2 were not taken into account for the sake of simplicity. Therefore, the main expected reactions during oxidation process are as follows:  O2 ðgÞ ! ZrO2 ðsÞ þ B2O3 ðlÞ  ZrB2 ðsÞ þ 5 2 LaB6 ðsÞ þ 21 O2 ðgÞ ! 1 B2O3 ðlÞ ! B2O3 ðgÞ 4 2 SiCðsÞ þ 3 O2 ðgÞ ! SiO2 ðlÞ þ COðgÞ 2  La2O3 ðsÞ þ 3B2O3 ðlÞ  ð1Þ  ð2Þ ð3Þ  ð4Þ  )  C  °  (  e  r  u  t  a  r  e p  m  e  T  2600  2400  2200  2000  1800  1600  1400  1200  1000  0  100  200  300 Time (s)  400  500  600  Figure 1. SEM micrograph of the polished surface of a ZrB2-20 vol.%  SiC-10 vol.% LaB6 sample.  Figure 3. Measured results of temperature curves of the ZrB2-20 vol.%  SiC-10 vol.% LaB6 sample.  Figure  2. Photographs  of  ZrB2-20 vol.% SiC-10 vol.% LaB6  (a)  during and (b) after oxyacetylene torch testing. The oxidized surface  Figure 4. Photograph and cross-sectional micrograph (the outer loose  scale has been removed) of oxidized ZrB2-20 vol.% SiC after  testing  layer appears compact and adherent, and no cracks or spallation were  under  the  same  conditions.  Spallation  and  cracks  (arrows) were  detected there.  observed after testing.    \\x0c', '1038  X.-h. Zhang et al.  / Scripta Materialia 57 (2007) 1036-1039  SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ SiO2 ðlÞ ! SiO2 ðgÞ La2O3 ðsÞ þ ZrO2 ðsÞ ! La2Zr2O7 ðsÞ  ð5Þ ð6Þ ð7Þ  The main reaction products from the oxidation of ZrB2 and LaB6 are zirconia, lanthana and amorphous boric acid (B2O3). The last compound has an unusually low melting point (450 °C) and a high vapor pressure (Figure 5). Therefore, at high temperature B2O3 quickly vaporizes (Reaction (3)) resulting in increased oxidation of ZrB2. The introduction of remarkably improves the oxidation resistance of the UHTC in the temperature ranging from 1200 °C to approximately 1700 °C. Nonetheless, at higher temperatures, SiC is no longer responsible for the remarkable improvement in oxidation resistance. On the one hand, SiC undergoes a transition from passive to active oxidation according to Reaction (4) at \\x181650 °C [9]. At \\x182400 °C, the oxidation resistance of SiC is poor due to active oxidation, which results in the formation of high vapor pressure oxide products like SiO and CO with almost no silica, which provides effective oxidation protection at elevated temperatures. On the other hand, the initially formed silica-enriched glass will be lost by the considerable evaporation due to the high vapor pressure of the SiO2 (g) (Figure 5) at very high temperatures in response to Reaction (6). The calculations were performed using JANAF data [10]. Furthermore, there was a large amount of H2O in the oxyacetylene torch testing, the molecule net of the silica being broken by forming non-bridging SiOH groups. This decreased the silica viscosity and increased the volatilization of SiO2 in the form of Si(OH)4, SiO(OH)2 and SiO(OH)). Therefore, silica is unstable in such severe environments. When the gases escape, they pass through the ZrO2, leaving behind channels of porosity. Because ZrO2 is not passivating and does not completely seal the surface of the sample, oxygen can diffuse through the cracks and porous oxide channels, and newly exposed UHTC surfaces are then subject to further oxidation. Another issue with ZrB2-SiC is the phase instability of its oxidation products. At high temperatures, ZrO2 is tetragonal. Upon cooling to room temperature they transform to the monoclinic structure with an attendant volume expansion. This phase transformation, coupled with their high thermal expansion coefficient and low thermal conductivity, can easily lead to cracking and spalling under thermal transient conditions.  Figures 6 and 7 show the SEM micrographs and EDS of the surface of oxidized ZrB2-SiC-LaB6 at 2400 °C for 600 s at low and high magnifications. The surface appeared compact, with only a few small microcracks and pores. EDS analysis of the surface of a oxidized sample shows that zirconium, oxygen and lanthanum are present as the primary elements, and silicon and boron are almost undetectable by the apparatus (Figs. 6 and 7). It is significant to note that the surface-oxidized sample consists of two distinctly different scales, namely, the bright and gray phases. The bright phase was distributed at the grain boundaries of the gray phase and sealed the cracks at the boundaries. X-ray diffraction results show the formation of m-ZrO2, t-ZrO2 and La2Zr2O7 as the major phases, as shown in Figure 8. The formation of t-ZrO2 was attributed to the solution of La2O3 in the ZrO2, which stabilized the ZrO2. Thus, the phase transformation of the ZrO2 on cooling to room temperature was impeded. The bright phase was identified as La2Zr2O7, according to the combination of the EDS and XRD results, which has a pyrochlore structure. The pyrochlores have ionic mobilities that are an order of magnitude lower than that of the high-temperature fluorite phase and have a melting temperature of >2300 °C [3]. The presence of La2Zr2O7 presumably derived from the chemical reaction between ZrO2 and La2O3 (Reaction (7)). The gray phase mainly consists of the ZrO2 and also contains some La2O3. The LaB6 addition overcomes the deficiencies of SiC-reinforced UHTC by promoting the formation of lanthanum zirconate, which has a higher melting point and viscosity, and efficaciously covers the sample surface and seals the cracks, effectively limiting the inward transport of oxygen. This results in  Figure 6. SEM micrograph and EDS of the surface of oxidized ZrB2- 20 vol.% SiC-10 vol.% LaB6 at 2400 °C for 600 s.  107 106 105 104 103 102 101 100 10-1 10-2 10-3 10-4 10-5 10-6 10-7 10-8 10-9  )  a  P  (  e  r  u  s s  e  r  p  r  o p a  V  10-10  1000  B2O3(g)  SiO2(g)  ZrO2(g)  1200  1400  1600  1800  2000  2200  2400  2600  Temperature (°C)  Figure 5. Vapor pressure vs.  temperature above some oxides, calcu Figure 7. SEM micrograph and EDS analysis of  the high magnifica lated at ambient pressure.  tion of the oxidized sample.      \\x0c', 'X.-h. Zhang et al.  / Scripta Materialia 57 (2007) 1036-1039  1039  Figure 8. XRD patterns of the oxidized surface.  an in situ, self-generating protective oxidation barrier and provides better oxidation protection up to 2400 °C. This suggests that the introduction of LaB6 provides tangible benefits for the resistance to oxidation at ultrahigh temperature. A cross-sectional image (a) and elemental maps for the distribution of (b) O, (c) B, (d) Si, (e) La and (f) Zr are shown in Figure 9. The oxide scale appears compact thickness of the scale is \\x18545 lm. No cracks or ruptures and adherent, with the exception of some voids. The total were observed in the scale after oxidized at 2400 °C for 600 s. Compositional analysis by EDS showed that the oxide scale was composed mainly of zirconium, oxygen and lanthanum. No oxide of B was detected in the whole oxide scale due to the loss of B2O3 by quick evaporation (Fig. 9(c)). Zirconium and oxygen are present along with lanthanum in the outside scale, which is free of silicon. This is consistent with the XRD results, which show the presence of ZrO2 and La2Zr2O7 at the surface. EDS map of silicon in Figure 9(d) shows a sharp transition at the boundary of the virgin material, corresponding  to an SiC-depleted region. The reduction of Si content in the innermost oxide layer is attributed to the active oxidation of SiC. Only minor amounts of the silicon were detected in the SiC-depleted layer. The loss of the silicon in the outer oxide scale is attributed to the volatilization of silica caused by the extremely high temperature and the permeation of H2O molecules through the scale. The distribution of the Zr element in both the unreacted material and the oxide scale is continuous and homogeneous, and no visible transition was detected, indicating that the oxide scale was coherent and compact even though a lot of voids were present. The content of La in the outer scale is richer than in the inner scale, as shown in Figure 9(e), which is consistent with the distribution of Si in the oxide scale of ZrB2-SiC at lower temperature (i.e. 1500 °C) [4]. The overall effect of the LaB6 addition is to significantly increase the capability to resist oxidative attack. On the one hand, La ions stabilize the tetragonal ZrO2, and the phase transformation accompanied with a volume expansion on cooling to room temperature is inhibited. On the other hand, the solution of La in the scale improves the melting point and viscosity of the oxides, and leads to the reduction of cracks and the formation of a coherent compact scale, which acts as a more efficient barrier against the inward diffusion of oxygen than SiO2 and ZrO2. In conclusion, the oxidation resistance of ZrB2-SiC- LaB6 composite was tested by oxyacetylene torch with a dwell time of 600 s and the surface temperature of the 2400 °C. The sample was over total weight 0.2% and \\x18545 lm, loss and thickness of the scale are only respectively. The addition of LaB6 significantly improves the oxidation resistance of ZrB2-SiC-based UHTC due to the formation of adherent compact scale without scale spallation.  This work was supported by the National Natural Science Foundation of China (90505015 and 50602010), the Research Fund for the Doctoral Program of Higher Education (20060213031) and the Program for New Century Excellent Talents in University.  J.  Eur.  [1] A.K. Kuriakose, J.L. Margrave, J. Electrochem. Soc. 111 (7) (1964) 827. [2] E.V. Clougherty, D. Kalish, E.T. Peters, AFML-TR-68190, 1968. [3] Kameleshwar Upadhya, Jenn-Ming Yang, P.Wesley Hoffman, Am. Ceram. Soc. Bull. 76 (12) (1997) 51. [4] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Ceram. Soc. 27 (6) (2007) 2495-2501. [5] A. Chamberlain, W.G. Fahrenholtz, G. Hilmas, D. Ellerby, Refrac. Applic. Transac. 1 (2) (2005) 1. [6] F. Monteverde, A. Bellosi, Adv. Eng. Mater. 6 (5) (2004) 331. [7] M.M. Opeka, UNITECR’05, in: Proceedings of the Unified International Technical Conference on Refractories 9th Biennial Worldwide Congress on Refractories, Wiley, New York, 2006, p. 522. [8] Gui-Ming Song, Yu Zhou, Yu-Jin Wang, Mater. Character. 50 (4-5) (2003) 293-303. [9] W.L. Vaughn, H.G. Maahs, J. Am. Ceram. Soc. 73 (6) (1990) 1540-1543. [10] M.W. Chase Jr., NIST-JANAF Thermochemical Tables, fourth ed., American Institute of Physics, Woodbury, NY, 1998.  Figure  9. A cross-sectional  image  (a)  and elemental maps  for  the  distribution of (b) O, (c) B,  (d) Si,  (e) La and (f) Zr.  \\x0c']"
},{
  "_id": 261,
  "PDF": "The effect of extreme temperature in an oxidising atmosphere.pdf",
  "Text": "['The effect of extreme temperature in an oxidising atmosphere  on dense tantalum carbide (TaC)  Anna Lashtabeg • Michael Smart  • Daniel Riley •  Andrew Gillen • John Drennan  Received: 14 March 2012 / Accepted: 11 July 2012 / Published online: 25 July 2012 Ó Springer Science+Business Media, LLC 2012  Abstract  This study describes the microstructure devel opment as dense tantalum carbide (TaC), which is subjected to extreme temperature environments (3,000 °C)  in  the presence of oxygen. These are conditions that structural  materials are expected to experience in hypersonic aero propulsion  applications. The  conditions  produce molten  oxide which may provide a temporary resistance to rapid  oxidation and may go some way to repair  thermal  shock  cracks, however, at  the same time the liquid is observed to  attack the dense  ceramic both chemically and mechani cally. A reaction mechanism is suggested which involves  dissolution of TaC in the oxide melt and a two step oxi dation;  first  the  reaction  of TaC with  oxygen  to  form  Ta(O,C) and TaOx,  resulting in dissolved dissociated car bon,  followed  by the  reaction of dissolved carbon with  oxygen to produce gas. This microstructural analysis of one  of  the candidate ultra-high temperature ceramic materials  for hypersonic flight provides new insight  into the mech anism of TaC oxidation and the role of  the liquid oxide  layer  in acting not only as  a protective  layer  to further  oxidation, as is commonly reported, but also as a dynamic  component that promotes erosion of the TaC surface and is  a source of  further oxygenation of  the TaC surface.  If  the  formation of  the liquid phase can be better controlled and  the reaction of  the liquid phase with the matrix be slowed  and stabilised,  then the formation of a liquid phase at  the  surface of TaC may provide a key to designing materials  that can withstand the rigours of hypersonic flight.  Introduction  Tantalum carbide (TaC)  is considered as a possible mate rial  for aero-propulsion use in extreme environments. The  developments in scramjet  technology that will propel aircraft at hypersonic speeds ([5 times the speed of sound, Mach number [5) will require the production of materials that can withstand temperatures in the vicinity of 3,000 °C  in an oxidising atmospheres. The search for such materials  has  concentrated  on  a  group  of  ultra-high  temperature  ceramics [1-5] of which TaC is one. With a reported melting temperature of 3,880 °C, this hard cubic material is  seen  as  a  potential material  to  be  used  as  a  protective  coating on leading edges in hypersonic vehicles. However,  in a similar way to other carbides and borides, TaC suffers  from poor oxidation resistance.  There  have  been  several  studies  on  the  oxidation  behaviour of TaC and related materials such as TaCN [5].  In each case,  they provide solid quantitative data on the  oxidation kinetics but in general the temperatures are at the  lower end of the spectrum. In particular Desmmaison-Brut  et al. [5] who has shown that a sample of dense TaC was oxidised fully at 850 °C in pure oxygen after 5 h producing (b-Ta2O5),  beta  tantalum pentoxide  the  low temperature  form of  this oxide. Opila et al.  [6]  suggested that  liquid  phase formation of possibly, Ta2O5 was observed after arc  jet  testing of some of  their composite materials; however,  this was only tentatively identified.  A. Lashtabeg \\x01 M. Smart  \\x01 J. Drennan (&)  The Centre for Microscopy and Microanalysis, The University of  Queensland, St Lucia, Brisbane, QLD 4072, Australia  e-mail:  j.drennan@uq.edu.au  A. Lashtabeg \\x01 M. Smart  \\x01 D. Riley \\x01 A. Gillen \\x01 J. Drennan  Defence Materials Technology Centre, Hawthorn, VIC 3123,  Australia  D. Riley \\x01 A. Gillen  Institute of Materials Engineering, Australian Nuclear Science  and Technology Organisation, Lucas Heights, NSW 2234,  Australia  123  J Mater Sci  (2013) 48:258-264  DOI 10.1007/s10853-012-6740-4  \\x0c', 'At extreme temperatures,  the metal carbides (MeC) and  metal  borides  (MeB) oxidation reaction at  the  solid/gas  interface produce  a  liquid oxide  layer  that  is  frequently  reported as protective in that  it stops further oxidation of  the solid carbide and boride. The nature of  the solid car bide/liquid  oxide  interface  has  never  been  previously  examined and the effects of the ‘‘protective’’ nature of the  liquid layer on the solid have not been studied.  A feature of the study reported in this article is that  the  fully dense ceramic is subject  to a designed oxy-acetylene  flame test  in an attempt  to simulate the extremes of  tem perature  and atmosphere  the material will  experience  in  hypersonic flight. Other dynamic high temperature systems  exist and have been used for the study of candidate carbide  and boride materials [7] but  to our knowledge, no studies  have  been  performed  on TaC. However,  some  of  the  microstructural features we describe have been observed in  high temperature  tests  and of particular  interest, porous  structures of surface oxides that often accompany oxidation.  In this article, we describe the interaction between dense  TaC and the liquid oxide layer  formed using an oxidising 3,160 °C. The  flame with  a  calculated  temperature  of  interaction  is  complex. The  oxidation  products  on  the  surface  become molten  during  the  test  and  this molten  phase  reacts with the bulk material  in a  series of  steps.  There  is  also  evidence  presented,  that  in  addition  to  chemical dissolution, mechanical erosion plays a part in the  process.  Experimental  Fully dense TaC was prepared by hot uniaxial pressing  TaC powder  (Atlantic Equipment Engineers, Grade TA 301, Bergenfield, NJ, USA)  and  subsequent  removal  of  remnant porosity by hot  isostatic pressing (HIP). Hot uni axal pressing of the loose TaC powder was conducted in a  graphite die, diameter 50 mm, using a resistively heated  element furnace (Elatec Inc., Andover MA, USA) operating at 2,000 °C for 1.5 h under vacuum (*0.1 MPa) and  an applied consolidating pressure of 35 MPa. The sample  was then hot  isostatically pressed (AIP Eagle 6, Columbus  OH, USA) using an argon atmosphere, a 1 h dwell time at a 2,000 °C  temperature  and  pressure  of  and  100 MPa,  respectively. The sample density was measured in accor dance with ASTM C20-2010, and found to be fully dense.  A test  specimen was  prepared  using  electric  discharge  machining (EDM), which produced a uniform sample disc  with dimensions of 2 mm thickness and 50 mm diameter.  Every effort was made to ensure that remnants of the EDM  process were removed before high temperature exposure.  This involved polishing the surfaces of the specimen. At no  stage  in the  characterisation was  any evidence of EDM  contamination  observed.  Preliminary  microstructural  characterisation revealed a fully dense specimen with an  average grain size of 15 lm.  Oxyacetylene flame tests were carried out on the Aus tralian Nuclear and Science and Technology Organisation  (ANSTO) ANSTO High Temperature Ablation Rig (Hi TAR)  facility, based at  the University of Sydney. The Hi TAR facility operates in general observation with ASTM  E285-08. The fully dense TaC specimen was loaded onto a  refractory sample holder  and subjected to gradual flame  heating using a variable-speed translation stage. The oxy acetylene torch tip was started 400 mm from the sample,  translating over  a 120 s period to within 19 mm of  the  surface, where it remained for a further 60 s. A calculated 3,160 °C was  adiabatic  flame  temperature  of  achieved  using an acetylene rich flame (fuel equivalence ratio 1.7)  with a flow velocity of 200 m/s. Using this test, the sample  surface temperature was measured using a dual wavelength  pyrometer (Mikron Quantum II, Oakland NJ, USA) to have reached a maximum of 2,100 °C. At  the completion of the  test,  the oxyacetylene torch fuel was  terminated and the  sample allowed to cool  to ambient.  After the flame test,  the sample was intact but consisted  of  the original  specimen, which was enveloped by a thin  slightly pink coloured material  that extended centimetres  from the surface of  the bulk of  the sample. This  surface  phase had spalled in the centre, while the remainder of the  surface appeared to be a clear glassy phase. The material  making up the thin, central over-layer was crushed to fine  powder  and subjected  to X-ray  powder  diffraction. The  envelope was  removed from the specimen and an X-ray  diffraction pattern recorded from the surface glassy region.  Samples for SEM analysis were sectioned, mounted in  conductive resin and polished to show the cross section from  the surface of  the sample through to the back side of  the  sample—the side that has not been in direct contact with the  flame. The polished specimen was examined in an analytical  scanning electron microscope  (JEOL 7001). Energy dis persive spectra (EDS) were recorded to confirm the presence  of oxygen and carbon in the grain boundary regions. How ever, as a result of the spot size limitation of the EDS system,  the analysis region was similar to the estimated interaction  zone of  the incident electron beam and as a consequence  could only reveal very general trends in composition.  To overcome the limited information from EDS, selec ted parts of  the polished specimens were examined using  micro-Raman  spectroscopy  (Renishaw  inVia  Raman  microscope)  to determine the presence and the nature of  any carbon dissolved in the system and to provide some  indication of other species present  in an attempt  to sort out  the complexity of the reaction. Using a 514 nm laser with a  509 objective  lens,  a maximum potential  resolution  of  1.5 lm may be achieved. In this study, large regions of the  J Mater Sci  (2013) 48:258-264  259  123  \\x0c', 'sample were examined to obtain a complete view of  the  various  chemical  trends  and  to minimise  any  selection  errors. All  recorded  spectra  showed  a  high  degree  of  reproducibility.  Results  X-ray  diffraction  results  from the  reacted  surface  are  complex, showing multiple high temperature phases, which  is  the  nature  of  tantalum based  oxides. This  is  further  complicated by the in-exact distribution of heat across the  sample  as a result of using a flame. Figure 1 presents a  series of diffraction patterns taken from different regions of  the heat  affected disc.  In Fig. 1a,  the diffraction pattern  shown is from the glassy coating at  the centre of  the disc,  which was exposed to the hottest part of the flame and was estimated to be close to 3,000 °C. It appears that  the phase  has been molten and the diffraction pattern could be pri marily indexed as multiple phases of  tantalum pentoxide.  The predominate phase was the high temperature form of  the oxide, a-Ta2O5, which has the monoclinic space group,  C121, with lesser quantities of  the low temperature poly morph, b-Ta2O5,  the orthorhombic form with space group  C2mm. The intensity anomalies  suggest  that a degree of  preferred orientation has occurred as the material has rap idly crystallised from the melt. The kinetics of  the trans formation from the high temperature a form to the  low  temperature b form is reported to be slow [8]. The condi tions of  this experiment appear  to have effectively quen ched the high temperature form of the oxide. It should also  be noted that the high temperature form can be quenched if  there are stabilisers present  [9].  In this experiment, where  violent  reactions  are  taking  place  on  the  surface,  it  is  possible that the high temperature phase has been stabilised  by other cations during this series of experiments,  the only  detectable ‘other cation’ was carbon.  The area of the disc away form the centre was exposed  to indirect flame, and the X-ray diffraction pattern (Fig. 1b)  shows  the  presence  a-Ta2O5 with b-Ta2O5. The peaks corresponding to a-Ta2O5 are broader  a  larger  quantity  of  and preferred orientation effects are not observed. Visual  inspection shows that  in this area of the sample,  the glassy  coating is thin and the Ta2O5 is primarily a powder  layer  seen peeling from the TaC surface. It  is estimated that  this  area of the sample reached the temperature of close to the (1,872 °C). The back of  melting point of  the Ta2O5  the  disk, which had no direct exposure to the flame,  showed  the primary phase  to be b-Ta2O5 small quantity of a-Ta2O5 a-Ta2O5 X-ray diffraction peak indicates that  and TaC, with only a  (Fig. 1c). An observed sharp  the presence  of some molten phase, which was later, confirmed to be the  case after examination in the SEM.  The  scanning  electron micrograph  shown  in  Fig. 2  reveals the interface between the molten tantalum oxide and  the remaining dense tantalum carbide in the central region  directly impacted by the flame. This is a cross section of the  specimen that  shows  the penetration of  the  liquid phase  from the surface of the sample. The images show the molten  oxide  liquid  phase  is  gradually  dissolving  the  tantalum  carbide and in Fig. 2,  the various stages of the dissolution  are  captured. Three  distinct  stages  are  observed:  grain  cleavage,  grain migration  through  a  liquid  phase,  grain  oxidation and dissolution. Grain cleavage is clearly seen  across the sample with boundaries increasing as liquid phase  infiltrates  into these grain boundaries. Various  stages of  grain cleavage are seen in Fig. 2. This was observed across  the sample in all areas, including the back of the sample that  did not experience the direct flame, thus indicating that  the  highly viscous liquid phase on the surface is responsible for  the attack. The mechanical erosion of TaC by the liquid  layer is clearly seen in the microstructure of the samples and  occurs in all areas at  the solid/liquid interface.  Grain migration through a liquid phase forms a distinct  layer  around the  sample,  showing islands of TaC grains  floating in the  liquid Ta2O5 phase. The  amount of TaC  grains  and the  thickness  and of  this  layer varies  across  different  sample regions.  In the area directly beneath the  flame the number of TaC islands is low, however  the area  extends further from the interface (200 lm) than in cooler  2θ (degrees)  20  30  40  50  60  70  (a)  (b)  (c)  *  *  *  *  *  *  *  *  *  *  *  *  *  *  *  *  *  *  *   Fig. 1 XRD data  for  different  regions  in  the  sample  showing  variation in phases  as a function of distance from the flame. a In the phase is a-Ta2O5 (asterisk). Anomalous  direct path of  the flame,  intensities are as a result of  texturing as the molten oxide solidified.  b Recorded from a region away from the direct flame. The main phase is a-Ta2O5. In this case the material has not been molten. c Recorded  from the back of the sample in the region shielded from the flame. Both TaC (open square) and b-Ta2O5 (filled circle) are observed with a minor amount of the a form which is thought to have leaked around  the sample from the molten region in the direct flame  260  J Mater Sci  (2013) 48:258-264  123  \\x0c', 'regions (50-100 lm). In the cooler area around the edges,  the  liquid  phase  is  seen  to  contain  significantly  higher  number of TaC islands, however  this area is smaller. The  back of  the pellet  that was not  in direct contact with the  flame shows  the smallest  ‘island’  interface. This  follows  the expected kinetics of migration and dissolution/oxida tion.  In the hottest area of the sample directly beneath the  flame, rapid oxidation of the TaC will occur, leaving fewer  grains  at  the  surface,  as  is  observed  in Fig. 2.  In  this  micrograph there is a clear gradation in TaC particle size  from the surface of  the specimen to the liquid phase/solid  TaC interface. The thickness of  this  region is a variable  depending on the distance from the direct flame. The back  of  the disk showed the smallest oxide layer thickness.  Dissolution is not instant and proceeds with a distinctive  front as can be seen in the micrographs. The depth of  the  penetration from the surface of  the specimen under  these  experimental conditions was up to 500 lm, about 20 % of  the original  sample dimension and this varied across  the  specimen depending on the temperature.  Across the sample, the point of attack of the liquid phase  is the grain boundaries and very little intra-granular dam age was  observed. The micrograph  shows  the  tantalum  oxide effectively separating out  individual grains of TaC,  which  then  reduce  in  size  the  further  away  from the  interface region.  In addition,  the porosity increases away  from the attack interface and as is shown later, is a result of  the evolution of trapped gas as the dissolved carbon forms  CO2 or CO. The attack is along grain boundaries as can be  seen in Fig. 2 and in more detail  in Fig. 3; however  in  some cases, as shown in the micrograph, (inset of Fig. 3),  attack  may  occur  directly  along  crystallographically  favoured  regions within  the TaC.  In Fig. 3,  inter-grain  cracking can be seen and this has been highlighted by the  attached arrows. It  is suggested that  this cracking maybe a  precursor to penetration by the liquid phase.  Energy  dispersive  spectra  (presented  in Fig. 4)  from  specific  regions,  reveals  the TaC grains which  do  not  contain any oxygen,  and an inter-granular phase  rich in  oxygen and containing carbon.  From these microstructural  observations,  the  liquid  oxide  layer  is  seen  to  attack  the  carbide  along  grain  boundaries. It  is possible that  the penetration of  the liquid  is  aided  by mechanical means,  similar  to  erosion  that  would be observed at a liquid/solid interface. The interface  region between the penetrating liquid and the dense TaC  would  be  subjected  to  considerable mechanical  stress,  resulting from the viscosity and flow of  the molten oxide  layer, which would  act  to weaken  the  grain  boundary  region.  It  is not necessary to evoke a mechanical compo nent to the mechanism of attack, however for completeness  we  suggest  that  the  cracking along the grain boundary,  observed  and  highlighted  in  Fig. 3,  is  evidence  of  mechanical  instability at  this  front. Both of  these mecha nisms would work in conjunction and the chemical attack  would be naturally exacerbated by the mechanical attack  mechanism, opening regions  for  infiltration and allowing  the  liquid  layer  to move  along  grain  boundaries  and  increase the surface area of the attack on individual grains.  The coupling of mechanical and chemical mechanisms is  seen in Fig. 3, where liquid phase ingress has been accel erated by strain before oxide infiltration.  The oxidation mechanism of metal carbides (MeC) have  been investigated previously and TaC oxidation is known  Fig. 2 A scanning electron micrograph showing the dissolution front  of the penetrating liquid oxide phase  Fig. 3  Scanning electron micrograph showing the dissolution front in  more detailed. Cracking proceeding the infiltration and separation of  the grains is indicated by arrows.  Inset  is a region where the attack  appears to be favoured along particular crystallographic directions  J Mater Sci  (2013) 48:258-264  261  123  \\x0c', 'to follow through the metastable Ta(O,C)x resulting in the  final formation of Ta2O5 and carbon [10-13]. In this study,  it  is proposed that  further steps to accommodate the vari able oxidation states of TaxOy, which in liquid form would  have high oxide mobility in the system and would be the  source of oxygen at  the TaC/Ta2O5 solid/liquid interface.  The  liquid  layer  chemically attacks  the  solid carbide  grain  through  oxidation  and  a  formation  of  TaCxOy  boundary as an intermediate structure.  In order  for  this to  happen, CO2 or CO evolution must occur and this produces  a porous structure in the oxide, all of which can be seen in  the  presented micrographs  (Fig. 4). Three macroscopic  regions  are distinguishable  in Fig. 4;  the TaC solid,  the  dense layer  immediately adjacent  to the TaC grains and a  central  region that appears  to have some porosity or par ticle inclusions. Clear boundaries exist between these lay ers  as  evidenced in the backscattered electron image of  Fig. 4, and a variety of analytical tools have shown there is  compositional variation across these layers.  It has been suggested that  the formation of  the oxide  layer may  act  as  a  protective  layer  against  further  oxidation;  however,  these  results  suggest  that  far  from  being a protective layer,  the oxide gradually dissolves the  carbide through a process of oxygen transfer. The nature of  the  interface between the  solid carbide  and liquid oxide  layers is complex in nature as evidenced by the micrograph  in Fig. 4 and although the molten oxide may act as a barrier  to oxygen diffusion,  the direct action of  the melt with the  solid TaC will be difficult  to prevent.  In oxide systems,  there are flexibilities  in the stoichi ometry of  the oxide that will allow for variable oxidation  reactions. Ta2O5 is readily reduced to Ta2O5-d, and at  the  point of  contact of  the  liquid with the TaC grains,  it  is  feasible  that  oxygen  transfer  between Ta-C and Ta-O  produces an intermediate zone between the two materials.  The thermodynamics of reaction between the TaC and the  Ta2O5 have been reported temperatures [2,500 °C,  previously  and  show that  at  the reaction is  spontaneous  [12]  with rapid reduction of the oxide and the formation of CO  and Ta, whereas  at  temperatures  \\x0c', 'the flame,  indicating the importance of diffusion and oxi dation kinetics at  this  interface. The non-porous  layer  is  distinct  from the porous  liquid layer,  indicating that CO2 evolution occurs at a distance of 0.5-5 lm from the TaC  grains. The  initial  oxidation  of  the TaC surface  from  available gaseous oxygen is rapid and the solid/gas inter face is rapidly replaced with the solid/liquid interface. The  formation of  the oxide layer  limits the availability of gas eous  oxygen  and  the  process  of  oxidation will  now be  primarily  through  anion  exchange  from the TaO to  the  TaC,  and  only  the  liquid  layer will  have  access  to  the  gaseous oxygen. The oxygen concentration will  thus vary  from the liquid/gas  interface to the liquid/solid interface.  Thus past  the  initial  step of  liquid layer  formation,  the  results indicate that  the rate of TaC and dissolved carbon  oxidation is diffusion limited. The porosity of  the liquid  Ta2O5  layer  increases with increasing distance  from the  TaC boundary, and again is a result of  the availability of  oxygen to oxidise dissolved TaC, TaCxOy and C species.  Towards the surface of the sample, the TaC grains continue  to dissolve, while  the porosity of  the  liquid phases dra matically increases as more of  the trapped carbon is oxi dised to gas (Fig. 5).  Raman  spectroscopy was  used  to  determine  if  free  carbon was  present  in  the  predominantly  oxide  phase  formed by the decomposition of TaC.  In the  spectra of  Fig. 6,  information was obtained,  in a progressive manner,  away from the interface of  the reacting liquid phase and  the solid TaC. Raman spectra were recorded at  the inter face, Fig. 6a, and 150 microns Fig. 6b and 300 microns  Fig. 6c  away from the  interface. The Raman spectra of  Fig. 6 have been corrected for background subtraction and 1,580 cm-1.  show two  distinctive  peaks  at  1,360  and  These peaks are definitive for disordered carbon [14] and  provide strong evidence for the presence of free carbon in  the melt. The  data  shows  that  the  amount  of dissolved  carbon decreases with increasing distance  from the TaC  interface. The Raman data confirms  the presence of  sig nificant quantities of dissolved carbon species in the mixed  oxide/carbide  layer  closest  to the TaC interface. At  the  outer edges of  the liquid phase, no evidence of dissolved  carbon species  is  seen. These results are consistent with  previous findings  and thermodynamic  equilibrium calcu lations, which show that in the TaC system (and other MeC  systems)  the metal  reacts first,  forming first  lower  then  higher oxides, followed by the oxidation of carbon to form  CO2 [10].  Current and previous work suggests  that  the TaC dis solves  in the liquid Ta2O5 and what may be a four  step  oxidation process via the formation of the species, TaCxOy  and C, followed by CO or CO2 evolution and the complete  tantalum oxidation. The  proposed  reaction  sequence  is  detailed  below and  the  important  reaction  sequence  is  numbered. The remaining reactions are simple final com plete oxidation of the liquid phase. The availability of free  oxygen may be the determining rate limiting step as  the  oxide layer will protect  the TaC from oxidation by atmo spheric oxygen. The  source of oxygen for  the oxidation  reaction of TaC is then the liquid tantalum oxide which in  turn is  reduced. This  reaction mechanism results  in the  complex microstructural variations we see.  Fig. 5 The region close to the surface of the sample where the TaC  grains are continuing to dissolve and where gas evolution from the  oxidation of carbon is producing a foam like microstructure  1000  1200  1400  1600  1800  2000  2200  2400  2600  2800  3000  Wavenumber cm-1  (c)  (b)  (a)  Fig. 6 Raman spectra taken at regular intervals from the TaC/Ta2O5 interface. a 150 lm, b 300 lm and c  close  to the  surface of  the  sample  J Mater Sci  (2013) 48:258-264  263  123  \\x0c', '1.  xTaC(s) ? 2.5xO2 ? TaC(s) ? xTa2O5(l) ? CO2(g) TaC(s) ? Ta2O5(l) ? TaCxOy(l) ? Ta2O5-d(l) ? C(l) TaCxOy(l) ? Ta2O5(l) ? 2Ta2O5-d(l) ? CO2(g) ? C(l) TaCxOy(l) ? O2(g) ? Ta2O5-d ? C  2.  3.  4.  TaCxOy(l) ? O2(g) ? Ta2O5(l) ? CO2(g) Ta2O5-d(l) ? d/2O2(g) ? Ta2O5(l) C(l) ? O2 ? CO2  The complexity and stoichiometric range of the TaCxOy  species  has  not  been  determined  but  possibilities  are  numerous. Raman  spectroscopy was  used  to  verify  the  presence of dissolved carbon species and supports previous  work on the oxidation of metal carbides. temperatures of 2,800-3,200 °C  The initial  reaction at  will  involve the formation of the Ta2O5 layer as a result of  the availability of oxygen at the interface between TaC and  air. With  the  formation  of  the Ta2O5  liquid  layer  the  availability  of  oxygen  at  the TaC interface  is  rapidly  diminished thus further  reactions are limited by the avail ability of air and must  therefore proceed through oxygen  exchange between the Ta2O5  liquid and TaC solid, as  is  shown  in  the  second  step. The  third  step  involves  the  reaction  of  the TaCxOy  species with Ta2O5  to  produce  Ta2O5, C and CO2. The final  reaction, occurring at  the  edges of the mixed layer, will  involve the oxidation of the  species in the liquid layer. From the Raman spectra, Fig. 6,  we observe decreasing carbon content as measurements are  made, away from the TaC/Ta2O5 interface.  Conclusion  These results show that at extreme temperatures and in the  presence of oxygen,  the oxidation of TaC and other car bides  is unexpectedly rapid. The resultant  formation of a  liquid oxide layer acts to slow the rate of oxidation of  the  carbide with gaseous oxygen. However,  the oxide layer  is  far  from being  protective  and  further  attacks TaC both  chemically and mechanically to degrade the structure at a  microstructural  level. At first sight,  this process appears to  be catastrophic, however better understanding of  this dis solution mechanism may provide  a means of protecting  these  non-oxide  surfaces  at  elevated  temperatures. The  formation of a molten oxide layer has advantages in that  it  can  uniformly  coat  a  surface,  repairing  the  inevitable  thermal  shock cracks, provide evaporative cooling and if  the kinetics of  the oxidation can be controlled, can effec tively slow the oxidation of  the main structural  compo nent—the non-oxide matrix. The possibilities are large in  that  these non-oxide systems can accommodate wide solid  solutions with  either  cations  or  anions  and  can  have  variable stoichiometry. In another approach, Shimada et al.  [15] showed more variable oxidation properties in powders  of TaC with changes in nitrogen content of the surrounding  gas. They found that the rate of oxidation is reduced by the  addition  of  nitrogen.  It  could  be  envisaged  that  if  the  composition of  the TaC was  slightly changed with addi tions of nitrogen then a more oxidising resistant material  could be made.  In addition,  the mechanical attack can be  minimised by changing the viscosity of the liquid layer and  the microstructure of  the  initial dense  ceramic.  In sum mary, understanding the complex nature of the carbide and  liquid oxide interface and the oxidation/erosion behaviour  of  the high temperature ceramic materials  is essential  in  designing future materials  for high temperature  applica tions.  The  above  observations  have  the  possibility  of  opening a new area in the search to solve the challenge of  protecting surfaces  in extreme environments by incorpo rating  active  surface  layers  that will  protect  all  the  important structural non-oxide material.  Acknowledgements  The  authors wish  to  acknowledge Mr Neil  Webb (ANSTO) for carrying out  the HIPing of the samples and Prof.  Assaad Masri  (University  of  Sydney)  for  access  to  the AMME  Combustion Laboratory for Hi-TAR testing to be  conducted. The  authors would also like to acknowledge the financial support of  the  Defence Materials Technology Centre (DMTC)  for  the funding and  directive of the research contained within this manuscript.  References  1. Wuchina E, Opila E, Opeka M, Fahrenholtz W, Talmy I  (2007)  Interface 16(4):30  2. Fahrenholtz WG, Hilmas GE, Chamberlain AL, Zimmermann  JW (2004)  J Mater  Sci  39(19):5951.  doi:10.1023/B:JMSC.  0000041691.41116.bf  3. Levine SR, Opila EJ, Halbig MC, Kiser JD, Singh M, Salem JA  (2002) J Eur Ceram Soc 22:2757  4. Carney CM (2009)  J Mater Sci 44:5673. doi:10.1007/s10853 009-3799-7  5. Desmaison-Brut M, Alexandre N, Desmaison  J  (1997)  J Eur  Ceram Soc 17:1325  6. Opila E, Levine S, Lorincz J  (2004)  J Mater Sci 39:5969. doi:  10.1023/B:JMSC.0000041693.32531.d1  7. Gasch M, Ellerby D, Irby E, Beckman S, Gusman M, Johnson S  (2004)  J Mater  Sci  39:5925.  doi:10.1023/B:JMSC.0000041  689.90456.af  8. Wu S, Chan H, Harmer MP (2005) J Am Ceram Soc 88(9):2369  9. Stephenson N, Roth RS (1971) J Solid State Chem 3:145  10. Shimada S (2002) Solid State Ion 149(3-4):319  11.  Johnsson M, Shimada S (2002) J Mater Sci Lett 21(12):955  12. Laurila T, Zeng K, Kivilahi JK (2002) Appl Phys Lett 80(6):938  13. Garg SP, Venkataramani R, Sundaram CV (1976) J Less Com mon Metals 50:245  14. Soukup L, Gregora I, Jastrabik L, Konakova A (1992) Mater Sci  Eng B11:355  15. Shimada S, Johnsson M, Urbonaite S (2004) Thermochim Acta  419:143  264  J Mater Sci  (2013) 48:258-264  123  \\x0c']"
},{
  "_id": 262,
  "PDF": "The effects of subsonic and supersonic dissociated air flow on the surface of ultra-high-temperature HfB2-30 vol_ SiC ceramics obtained using the solgel method.pdf",
  "Text": "['Journal of the European Ceramic Society 40 (2020) 1093-1102  Contents lists available at ScienceDirect  Journal of the European Ceramic Society  jou rna l homepage : www .e lsev ie r .com / loca te / jeu rce ramso c  Original Article  The effects of subsonic and supersonic dissociated air flow on the surface of ultra-high-temperature HfB2-30 vol% SiC ceramics obtained using the solgel method  Elizaveta P. Simonenkoa,*, Nikolay P. Simonenkoa, Andrey N. Gordeevb, Anatoly F. Kolesnikovb, Anton S. Lysenkovc, Ilya A. Nagornova,d, Vladimir G. Sevastyanova, Nikolay T. Kuznetsova  a Kurnakov Institute of General and Inorganic Chemistry of the Russian Academy of Sciences, Leninsky pr. 31, Moscow, 119991, Russia b Ishlinskii Institute of Problems of Mechanics of the Russian Academy of Sciences, 101-1 pr. Vernadskogo, Moscow, 119526, Russia c A.A.Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Leninskii pr. 49, Moscow, 119334, Russia d Dmitry Mendeleev University of Chemical Technology of Russia, 9 Miusskaya sq., Moscow, 125047, Russia  T  A R T I C L E  I N F O  A B S T R A C T  Keywords: Sol-gel processes Nanomaterial Ceramics UHTC Induction plasmatron Arc-jet testing High-enthalpy air flow  By hot pressing (1900 °C, 30 MPa, holding time 15 min) of HfB2-(SiO2-C) composite powder, an ultra-hightemperature ceramic material of HfB2-30 vol% SiC composition with nanocrystalline silicon carbide has been obtained. The effects of subsonic and supersonic dissociated air flow on the surface of produced materials have been studied on a high-frequency induction plasmatron in the geometry of a cylindrical sample with a flat face, fixed in a copper water-cooled holder with a 1 mm protrusion. It has been shown that a sudden rise in the average surface temperature of the samples to 2600 °C is characteristic for both modes, which is associated with the occurrence of local sites with a temperature > 2000 °C and a subsequent increase in their area. It is a matter of evaporation of the borosilicate glass layer from the surface of the oxidized sample and formation of the ceramic layer of a highly catalytic and low thermal conductive porous HfO2 layer, which is confirmed by the emission spectroscopy data, XRD and elemental analysis of the material surface after the experiments. The features of heating the oxidized surface of the samples under the impact of subsonic and supersonic dissociated air flow have been noted: there are differences in the location of overheated sites, initiating a sharp temperature rise and the rate of growth of their area.  1.  Introduction  Scientists and technologists from different countries are interested in developing materials that can withstand aerodynamic heating to temperatures above 2000-2500 °C: in particular, materials that are based on zirconium/hafnium diboride and silicon carbide [1-18]. The use of such ultra-high-temperature ceramics (UHTC) is promising not only for creating parts of hypersonic aircraft and propulsion systems, but in recent works, it has also been proposed to use materials based on refractory metal borides, such as solar absorbers [19-21]. This is due to the unique combination of such qualities as high melting points, thermal conductivity (including at temperatures up to 2000 °C), oxidative stability (including in high-speed air flow), good mechanical properties for ceramics and optical characteristics suitable for the above  applications. The use of highly dispersed ZrB2/HfB2 and SiC initial powders when producing UHTC, as shown in [22-29], makes it possible to stimulate the consolidation process and significantly improve the mechanical properties of the materials. In addition, it was noted in [29] that when replacing SiC powder with an average particle size of 44 μm by nanoand submicron powders for samples of ZrB2-20 wt.% SiC ceramics obtained by hot pressing, a significant increase in oxidative stability can be observed at temperatures of 1500 and 1700 °C. This fact is especially significant, since oxidative stability in the presence of oxygen at elevated temperatures is one of the key requirements for UHTC. The sol-gel method is extremely useful for the relatively low-temperature synthesis of nanodispersed refractory carbides [30-34], including silicon carbide [35-39]. It also has significant advantages for  ⁎ Corresponding author. E-mail addresses: ep_simonenko@mail.ru (E.P. Simonenko), n_simonenko@mail.ru (N.P. Simonenko), a_gord@mail.ru (A.N. Gordeev), koles@ipmnet.ru (A.F. Kolesnikov), toxa55@bk.ru (A.S. Lysenkov), il.nagornov.chem@gmail.com (I.A. Nagornov), vg_sevastyanov@mail.ru (V.G. Sevastyanov), ntkuz@igic.ras.ru (N.T. Kuznetsov).  https://doi.org/10.1016/j.jeurceramsoc.2019.11.023 Received 5 July 2019; Received in revised form 4 November 2019; Accepted 6 November 2019 Available online 13 November 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.  \\x0c', 'E.P. Simonenko, et al.  the production of highly dispersed ZrB2(HfB2)-SiC composite powders [40-45], with their subsequent use for the production of ceramics or the application of protective antioxidant coatings on the surface of carbon fiber-reinforced ceramic-matrix composites. In our previous works [3,8,46,47], it was proposed to combine the stages of nanodispersed silicon carbide synthesis and UHTC production at moderate temperatures of 1700-1900 °C, i.e., to carry out hightemperature compaction of HfB2-(SiO2-C) obtained by the sol-gel method, in accordance with the ideology of reaction sintering. The UHTC of HfB2-xSiC composition (x = 10-65 vol%) obtained in a similar way were studied for resistance to airflow oxidation in the thermal analysis mode at temperatures of 20-1400 °C [46]. And for HfB2-30 vol % SiC materials synthesized at minimum temperatures of 1700 and 1800 °C, single tests were carried out under the influence of subsonic [8] and supersonic dissociated air flow on a high-frequency induction plasmatron [3,47], which showed that they behave at least as well as those obtained during spark plasma sintering of coarse-crystalline SiC and HfB2 powders [9-11]. The objective of this work is to compare the behaviour of ultra-hightemperature HfB2-30 vol% SiC ceramic composites obtained by hot pressing of a highly dispersed HfB2-(SiO2-C) composite powder at a higher temperature of 1900 °C, under the influence of subsonic and supersonic dissociated air flow of a high-frequency induction plasmatron.  2. Materials and equipment  Reagents used: Tetraethoxysilane (TEOS) Si(OC2H5)4 (> 99.99 %, EKOS-1 JSC), LBS-1 bakelite varnish (Karbolit OJSC), formic acid CH2O2 (> 99 %, Spektr-Chem LLC), hafnium diboride (> 98 %, particle size 2-3 microns, aggregate size  20-60 microns, Tugoplavkie Materialy LLC). Production of ultra-high-temperature HfB2-30 vol% SiC ceramics was carried out using a hot press manufactured by Thermal Technology Inc. (model HP20-3560-20) [3,8,46,47]. HfB2-(SiO2-C) and composite powder was obtained by the sol-gel method [41,46,47]. Hot pressing was carried out at a temperature of 1900 °C (heating rate 10°/min) and a pressure of 30 MPa; holding time at a given temperature was 15 min. A small amount of boron nitride was used as a mould lubricant. The X-ray diffraction patterns of the synthesized composite powders were recorded on a D8 Advance (Bruker) X-ray diffractometer in the range of 2Θ 34‒37° with a resolution of 0.02°, the signal being accumulated for 2 s at the point, and in the range of 2Θ 5-80° with a resolution of 0.02°, the signal being accumulated for 0.3 s at the point. The infrared reflectance spectra of ceramic materials were recorded using an InfraLUM FT-08 FTIR spectrometer (PIKE EasiDiff diffuse reflectance accessory). Scanning electron microscopy (SEM) data were obtained with a triple-beam workstation NVision 40 (Carl Zeiss); the elemental composition of microdomains was determined with an EDX system (Oxford Instruments). The thermal behaviour of the products in the air flow (100 mL/min) in the temperature range of 20-1400 °C (heating rate 20°/min) was studied using the combined TGA/DSC/DTA analyzer SDT Q-600. Experiments on the effects on samples of a high-enthalpy jet of dissociated air were performed on the 100-kilowatt VGU-4 induction highfrequency plasmatron. Both for subsonic and supersonic modes, a steplike increase in the anode power of the plasmatron (N) from 20 to 70 kW in increments of 10 kW was made; the holding time at each step was three minutes. After reaching N = 70 kW, the sample was held at this power until the experiment was completed, and the total holding time was 40 min. The sonic nozzle with an output section diameter of 30 mm was used both in subsonic and supersonic modes. The experimental conditions for two different modes are given in more detail in Table 1. The sample surface temperature was measured using a Mikron M-770S pyrometer in the spectral ratio pyrometer mode (the temperature range being 1000-3000 °C, the diameter of the viewing area  Journal of the European Ceramic Society 40 (2020) 1093-1102  Table 1 Modes of exposure of the sample surface to subsonic and supersonic dissociated air flows.  Parameter  Sample No. 1 Subsonic flow  Sample No. 2 Supersonic flow  Air flow rate, g/s Distance between the exit from the sonic nozzle and the sample, mm Pressure in the plasmatron chamber, hPa Heat fluxb q, W/cm2  2.4 30  3.6 25  100 163-574c  16-20a 232-779c  a  In a supersonic mode, such a change in pressure in a submerged space does not affect the realized heat fluxes and stagnation pressure. b The heat flux was measured by a water-cooled copper calorimeter. c Changed depending on the anode power of the plasmatron.  being 5 mm). To record the temperature distribution over the sample surface, the Tandem VS-415U thermovisor was used: the measurements were carried out with the spectral emissivity set at ε = 0.6 at a wavelength of 0.9 μm [3,8-11,47]. This value corresponds to the emissivity of the oxidized layer based on porous HfO2. This emissivity was determined in previous comparative experiments on the simultaneous determination of the surface temperature of similar samples using the Tandem VS-415U thermovisor and the Mikron M-770S spectral ratio pyrometer in such a way, that the readings of both instruments are equal. At surface temperatures < 1800-1900 °C (when transparent silicate glass prevailed on the surface), the real value of ε was 0.75-0.8, as a result of which the error in determining the surface temperature in the range of 1200−1900 °C was  70−110 °C. The stagnation pressure Pst was measured by a water-cooled Pitot tube with a hemispherical toe of R =15 mm and a receiving opening diameter of 14 mm. The method for recording the emission spectra of the boundary layer above the sample surface is described in detail in [10,11,47]. For this purpose, the HR-4000 high-resolution compact diffraction spectrometer (Ocean Optics, USA) with a linear CCD detector (3648 pixels) with fiber-optic radiation input was used. In this experiment, a spectral range of 200÷650 nm was recorded.  3. Results and discussion  3.1. Production and study of HfB2-30 vol% SiC ceramics  Samples of ultra-high-temperature HfB2-30 vol% SiC ceramics were obtained as a result of hot pressing of the HfB2-(SiO2-C) composite powder synthesized by the sol-gel method. Synthesis of the initial composite powder was carried out during controlled hydrolysis of tetraethoxysilane (acid catalysis with formic acid) in the presence of a polymer carbon source (phenol-formaldehyde resin LBS-1, which turns into amorphous carbon after subsequent pyrolysis) [36,39] and dispersed HfB2 powder [3,8,41,47]. After gelation (due to which the maximum uniform distribution of all components in the volume was recorded), drying and carbonization at the temperature of 400 °C under dynamic vacuum conditions (residual pressure 1-5·10-6 atm), a product was produced, in which the most reactive SiO2-C system was formed on micron particles of hafnium diboride. The high dispersion ability of both the components of this system and the high uniformity of their distribution relative to each other, partially inherited from a true solution of phenol-formaldehyde resin and tetraethoxysilane, allows for the carbothermic synthesis of nanocrystalline silicon carbide during hot pressing at relatively low temperatures. In this experiment, the hot pressing temperature was limited to 1900 °C, with a holding time of 15 min. As the result, cylindrical samples were obtained with a diameter of 15 mm and a thickness of 4.3-4.5 mm. The density of the produced samples was 7.8 ± 0.3 g/cm3, which is 89 % of the calculated value obtained by the additive method (HfB2  1094  \\x0c', 'E.P. Simonenko, et al.  Journal of the European Ceramic Society 40 (2020) 1093-1102  Fig. 1. X-ray diffraction pattern of the original HfB2-(SiO2-C) composite powder (1) and the produced HfB2-30 vol% SiC ceramics (2); the inset shows a larger area of 2θ = 35-37°, in which the most intense reflex of the silicon carbide phase occurs.  density being assumed to be 11.2 g/cm3 [48], and SiC density - 3.2 g/ cm3 [49]). According to the XRD data (Fig. 1), as the result of hot pressing, nanocrystalline silicon carbide was synthesized; the reflexes for the cubic silicon carbide phase [51], which have a much lower intensity, were added to those of the HfB2 phase [50] on X-ray patterns. No traces of by-product phases (SiO2 [52], HfO2 [53] and HfC [54]) were found. The average size of SiC crystallites estimated by the Scherrer method was 38 ± 4 nm. Synthesis of silicon carbide from SiO2-C was also confirmed by the data of IR spectroscopy of diffusion reflection: in the range of wave numbers 800-950 cm-1, the absorption band characteristic of ν(Si-C) appeared, while the broadened absorption band at 970-1220 cm-1 associated with stretching vibrations of the Si—O groups disappeared. Scanning electron microscopy showed (Fig. 2) that the distribution of HfB2 particles and synthesized SiC was fairly uniform. The average size of hafnium diboride grains was 1-3 μm, i.e., the use of this method makes it possible to destroy the initial HfB2 aggregates, the size of which reached 60 μm. Thermal analysis of the obtained UHTC behaviour in air flow at temperatures of 20-1400 °C (Fig. 3) indicates that intense oxidation of the sample (of near-surface areas, first of all), accompanied by an increase in mass, begins at temperatures > 670-700 °C. This process corresponds to an exothermic effect at 740 °C. Some change in the process kinetics at higher temperatures is substantiated by the complication of the oxidation mechanism, taking into account the stage of diffusion through the borosilicate glass layer formed on the surface. At a temperature of 1298 °C, the maximum mass gain of 0.17 % was  Fig. 3. DSC (blue) and TGA (red) curves of the produced HfB2-30 vol% SiC material. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)  recorded. The tendency towards a decrease in the sample mass that was revealed at a higher temperature can be explained by an increase in the intensity of evaporation processes for the vitreous layer components on the surface, primarily of boron oxide. In general, it is possible to state the high oxidative stability of the obtained samples when they are heated in air flow (250 mL/min) to the temperature of 1400 °C. The compression strength of the obtained HfB2-30 vol% SiC composite with the porosity of 11 % exceeded 555 MPa, which corresponds to the data obtained in [46].  3.2. Studying the behaviour of ultra-high-temperature HfB2-30 vol% SiC ceramics under the impact of subsonic and supersonic dissociated air flow  In order to study the behaviour of HfB2-30 vol% SiC UHTC under the impact of high-enthalpy dissociated air flows, the samples were fixed in a composite copper model [3,8,47]. The mandrel into which the sample was inserted was pushed via a sliding fit into a water-cooled holder. The thermal contact of the face surface of the model parts was ensured by springing of the tension pin; the contact surfaces were lubricated with thermal grease to improve heat transfer from the mandrel to the water-cooled holder. Samples were installed in a water-cooled model with a protrusion of 1 mm relative to the front surface [10,11,47] using three cotton yarns based on SiC whiskers, which prevented direct contact of the sample with the copper mandrel and reduced heat transfer. Samples were introduced into a jet of dissociated air at the anode power of the plasmatron of 20 kW, which gradually increased to 70 kW in increments of 10 kW. The specific values of the exposure parameters and the average temperature of the sample surface determined using a spectral-ratio pyrometer are given in Table 2 as well as in Fig. 4. The behaviour of the samples under the described conditions is considered in more detail below.  Fig. 2. The microstructure of the produced HfB2-30 vol% SiC ceramic and distribution of the silicon and hafnium elements.  1095  \\x0c', '3.2.1. Sample No. 1, behaviour in subsonic dissociated air flow As can be seen in Fig. 4a, a corresponding increase in the average temperature of the sample surface occurred with a gradual increase in the anode power of the plasmatron. However, at the fifth stage (N = 60 kW, 12-15 min), there appears to be a certain tendency for the average temperature to rise within 3 min by 75 °C, and when switching to the anode power of 70 kW (q = 574 W/cm2), a sharp temperature rise of the sample surface occurs, which is characteristic of such materials and was noted both in our previous works [8,10,11] and in [55,56]. The surface temperature rise lasts up to the 25th minute for 10 min, when the value of 2610-2630°C is stabilized. The processes taking place were studied in more detail by analyzing thermal images at different points of exposure - Fig. 5. As can be seen, at the initial stages of exposure (N = 20-40 kW), the temperature distribution over the surface is, in general, uniform; however, when the anode power supply is 50 kW (q = 425 W/cm2), small areas appear at the sample edges with a temperature exceeding the average one (1540-1580 °C) by 30-50 °C (indicated by arrows in Fig. 5). The next increase in N to 60 kW leads to an increase in this difference: at an average surface temperature of 1650 °C, the temperature in local ‘overheated’ areas is 1750-1780 °C. The transition to power N = 70 kW (q = 574 W/cm2) is accompanied by a sharp increase in both the temperature and the number of such areas (mainly at the sample periphery): at the 16th minute of exposure with an average surface temperature of about 1700 °C, their temperature reaches 1900-1950 °C and quickly rises to 2300-2650 °C. As a result of increasing the area of hightemperature sites up to the 25th minute of the experiment, the average surface temperature rises to 2600-2700 °C, and the minimum temperature exceeds 2400 °C. In this case, there were points (Fig. 6) at  \\x0c', 'E.P. Simonenko, et al.  Journal of the European Ceramic Society 40 (2020) 1093-1102  Fig. 5. Thermal images of the surface of sample No. 1 of the HfB2-30 vol% SiC composition at various moments of exposure to a subsonic dissociated air flow; images corresponding to a sharp increase in the average surface temperature are highlighted in pink, the stage of sample cooling is highlighted in blue. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)  the HfB2-30 vol% SiC sample is associated with the appearance of local areas with elevated temperature at some moment and an increase in their area (Fig. 9). However, a number of peculiarities can be noted for sample No. 2. Thus, under the impact of a supersonic flow, areas with a temperature of > 2000 °C are formed not at the sample periphery (as we noted for sample No. 1 and in [8,10,11], as well as in Marschall et al. [55]), but in the central area. Such specific features of heating the HfB2-SiC UHTC surface by supersonic air flow were noted in [3,47]. In this case, at the 15th minute of the present experiment, the formation of a quite significant (with a diameter of 5 mm) area with elevated temperature (highlighted with an ellipse in Fig. 9) was clearly observed near the centre of the sample: at an average surface temperature of 1580-1610 °C, its temperature was 1630-1645 °C (Fig. 10). Approximately 15 min and 15 s after the start of the experiment, small (with a diameter of 0.2-0.5 mm) areas began to appear in this area with an even higher temperature: 50-100 °C higher, than the average temperature, which in just one minute heated to the temperature > 2000-2100 °C (see Fig. 10) and increased their area. It should be noted, that from 16 min 39 s and further, a sharp rise in temperature to 2750-2800 °C was observed due to the formation of protuberances on the sample face (Figs. 9-11), which may be caused by the release of borosilicate glass from the inner areas of the sample, with its further intense evaporation due to extremely rapid heating. At the end of the 17th minute, the temperature of the entire surface of sample No. 2 exceeded 2500 °C, and only the temperature at the very edge of the sample (ring 1-1.5 mm wide) did not exceed 2200-2300 °C (Fig. 11). Fig. 11 also clearly shows that, unlike sample No. 1, the difference in surface temperature of sample No. 2 by 100-130 °C in different sites of the sample (on the edge or in the centre) was peculiar to a supersonic impact right at the beginning of the experiment (N = 20 kW, q = 232 W/cm2). This situation was repeated for all subsequent stages up to N = 60 kW (q = 691 W/cm2), when active evaporation of borosilicate glass from the surface and exposure of the porous HfO2 layer began, which caused the appearance of areas with a temperature of > 2000 °C. After stabilization of the surface temperature values, the difference between the areas in the centre (on protrusions - point 1, Fig. 11) and on the edge (point 2, Fig. 11) was 450 °C. When the heat was turned off after a 40-minute exposure  to a  Fig. 6. Temperature variation at different points on the surface of sample No. 1 of the HfB2-30 vol% SiC composition when exposed to a subsonic dissociated air flow.  3.2.2. Sample No. 2, behaviour in supersonic dissociated air flow In Fig. 4b, it can be seen that the behaviour of sample No. 2 under the impact of supersonic flow during the first 12 min of exposure (N = 20-50 kW) is similar to that for sample No. 1 (subsonic flow) - with a gradual increase in the anode power supply of the plasmatron and, accordingly, in the heat flux, the average surface temperature increases from 1250 to 1620 °C (according to a spectral-ratio pyrometer). An increase in N to 60 kW (q = 691 W/cm2) in the first minute of exposure leads to a corresponding increase in the average temperature, but at the beginning of the 14th minute the tendency changes: instead of stabilizing the temperature with a slight decrease, it increases by 55 °C in 1.5 min: from 1718 to 1773 °C. A further increase in the anode power supply to 70 kW led to a very rapid increase in the average surface temperature to 2700 °C within 2 min. The subsequent slight decrease in temperature led to stabilization of an average surface temperature of 2685- 2695 °C at the 21 st minute, which did not change until the heating was turned off. As in the case of sample No. 1, under the impact of supersonic dissociated air flow, a sharp rise in the average surface temperature of  1097  \\x0c', 'E.P. Simonenko, et al.  Journal of the European Ceramic Society 40 (2020) 1093-1102  Fig. 7. Temperature distribution over the diameter of sample No. 1 at various moments in the experiment (see Fig. 5).  The sharp decline of the surface temperature to 1000−1500 °C (at the rate 350-600°/s for different surface areas) can be considered as a variant of the hardening material. That is, the microstructure observed for the samples, the elemental and phase compositions of the surface can be related with some approximation to those in the last seconds of exposure to highly enthalpy air flows. As can be seen, there are real protrusions on the oxidized surface, for which, judging by thermal images (Figs. 5-7 and Figs. 9-11), the highest temperatures were characteristic. The protrusions for sample No. 2 very much resemble residues of bursting large bubbles (with the diameter of 3-5 mm). X-ray analysis of the oxidized surface of the samples showed (Fig. 13), that monoclinic hafnium dioxide is the only crystalline phase [53]. This is due to the fact that the surface temperature during exposure, even in the least heated areas, significantly exceeded 1850-2000 °C in both experiments. This led to the evaporation of boron and silicon oxides from the surface that makes up borosilicate glass, which was produced during the oxidation of HfB2-SiC and performed protective functions at lower temperatures (< 1500-1700 °C). More refractory hafnium oxide with significantly lower vapour pressure occurred on the sample surface. Thus, according to [59], the partial pressure of HfO vapour above HfO2 at the temperature of 2108 °C is 8.44·10-5 mm Hg, while at a much lower temperature of 1315 °C, the pressure of B2O3 vapour above the SiO2-B2O3 system is in the range of 0.10-1.58·10-3 mm Hg [60]. According to the SEM data (Fig. 14), the surface microstructure of HfB2-30 vol% SiC UHTC samples after 40 min’ exposure to subsonic and supersonic dissociated air flows is, in general, similar - a porous ceramic layer is formed. However, it should be noted, that the oxidized surface layer is looser for sample No. 1, which is characterized not only by regular-shaped pores with a diameter of 1-7 μm but also by larger pores with a diameter of up to 100 μm with chaotic edges. Study of the atomic number contrast mode confirmed that there are no phase inclusions besides HfO2. It should be noted, that there are cracks in the surface HfO2 layer (Fig. 14b, c, e, f) associated with polymorphic transformations of this substance and a significant change in volume upon cooling. Cracks advantageously distributed between the grains and decelerated at  Fig. 8. The emission spectrum (wavelength intervals of 246-254 and 286-292 nm)1 of the boundary layer above the surface of sample No. 1 of the HfB2-30 vol% SiC composition; 17 min 16 s.  supersonic dissociated air flow, the surface temperature decreased by 1000-1370 °C (Figs. 9, 11) within 3 s; however, no cracking or detachment of the oxidized sites was observed. This is probably due to the braking of cracks on micropores on the surface of HfO2. The total mass loss of the sample was 3 %.  3.3. Studying the composition and microstructure of ultra-high-temperature HfB2-30 vol% SiC ceramics after exposure to subsonic and supersonic dissociated air flow  Fig. 12 shows the appearance of the samples after the experiments.  1 According to [57,58], the sensitive lines of boron are 249.678 and 249.773 nm, of silicon - 250.690, 251.432, 251.611, 251.921, 252.412, 252.851, and 288.158 nm  1098  \\x0c', 'E.P. Simonenko, et al.  Journal of the European Ceramic Society 40 (2020) 1093-1102  Fig. 9. Thermal images of the surface of sample No. 2 of the HfB2-30 vol% SiC composition at various moments of exposure to a supersonic dissociated air flow; images corresponding to a sharp increase in the average surface temperature are highlighted in pink, the stage of sample cooling is highlighted in blue. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)  present micropores. The study of the elemental composition of the surface by the EDX analysis made it possible to confirm that hafnium dioxide is on the surface. In this case, the n(Hf):n(Si) ratios for a material exposed to supersonic air flow (sample No. 1) are more, than 3.5 times higher than those for sample No. 2 - Table 3. For sample 2, the n(Hf):n(Si) ratio near the bursting bubbles was additionally studied. It has been established, that the amount of hafnium inside the crater is 150 times higher than that of silicon, and the n(Hf):n(Si) = 32 ratio (i.e., the silicon content) on the bursting bubble shell, which forms a certain annular  protrusion,  is elevated if compared to the average one on the surface.  4. Conclusions  As part of the hybrid method, which includes the sol-gel synthesis of HfB2-(SiO2-C) composite powder and its hot pressing under moderate conditions (1900 °C, 30 MPa, holding time 15 min), combined with the carbothermic synthesis of nanocrystalline silicon carbide, cylindrical samples of ultra-high-temperature HfB2-30 vol% SiC ceramics were produced. The density of the obtained samples was 7.8 ± 0.3 g/cm3,  Fig. 10. Temperature distribution along the diameter of sample No. 2 of the HfB2-30 vol% SiC composition for 15-17 min of the experiment (jump in the average surface temperature): the appearance of local overheated areas (a) and an increase in their area (b) (see Fig. 9).  1099  \\x0c', 'elevated temperature (Figs. 5, 9). However, for sample No. 1, a typical  which is 89 % of the calculated value. The produced UHTC samples of the same composition and structure were subjected to gradual heating with dissociated air flows: sample No. 1 with a subsonic flow and sample No. 2 with a supersonic flow. It should be noted, that the obtained materials, in which silicon carbide is nanocrystalline and, accordingly, highly chemically active, nevertheless survived the long-term (40-min) exposure to high-speed air flows, consisting of fully atomized oxygen and partial nitrogen, without destruction and complete oxidation. In this case, it should be noted, that there are some differences in the course of the surface oxidation of these materials by subsonic and supersonic air flows. For both experiments, then, there is a jump in the average surface temperature to values of 2500-2700 °C (according to a spectral-ratio pyrometer) after some long enough (> 12-14 min) exposures, but the time required to stabilize the maximum average temperature varies greatly. For sample No. 1, which was exposed to subsonic air flow, the time from the onset of temperature rise until reaching the plateau was 9-10 min, and for sample No. 2, this process occurred within only 2-3 min (Fig. 4). As the result of analysis of thermal images of the sample surface during heating, it was established, that the increase in the average surface temperature is due to a gradual increase in the area of sites with  \\x0c', 'E.P. Simonenko, et al.  Journal of the European Ceramic Society 40 (2020) 1093-1102  Fig. 14. The microstructure of the surface of HfB2-30 vol% SiC after exposure to subsonic (a-c) and supersonic (d-f) dissociated air flows; (a, b, d, e) - according to the data of the secondary electron detector (c, f) - in the average atomic number contrast mode.  Table 3 The n(Hf):n(Si) ratio on the surface of HfB2-30 vol% SiC UHTC samples after exposure to subsonic and supersonic dissociated air flows, EDX.  Sample name  Sample No. 1 (subsonic flow) Sample No. 1 (supersonic flow)  n(Hf): n(Si)  14 50  high-temperature HfB2-30 vol% SiC ceramics with nanocrystalline silicon carbide for use under long-term exposure to high-speed dissociated air flows (both subsonic and supersonic) at the temperatures above 2500 °C.  Declaration of Competing Interest  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.  Acknowledgements  The study has been carried out with financial support from the grant of the Russian Science Foundation (No. 17-73-20181).  References  [1]  [2]  [3]  [4]  E.P. Simonenko, D.V. Sevast’yanov, N.P. Simonenko, V.G. Sevast’yanov, N.T. Kuznetsov, Promising ultra-high-temperature ceramic materials for aerospace applications, Russ. J. Inorg. Chem. 58 (2013) 1669-1693, https://doi.org/10.1134/ S0036023613140039. E.P. Simonenko, N.P. Simonenko, V.G. Sevastyanov, N.T. 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  "_id": 263,
  "PDF": "The Oxidation Behavior of ZrB2-Based Mixed Boride.pdf",
  "Text": "['Key Engineering Materials Vol. 403 (2009) pp 253-255 © (2009) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.403.253  Online: 2008-12-15  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, www.ttp.net. (ID: 142.103.160.110, University of British Columbia, Kelowna, Canada-07/07/15,14:12:35)  The Oxidation Behavior of ZrB2-based Mixed Boride Seung Jun Leea, Do Kyung Kimb  Department of Materials Science and Engineering, KAIST, 373-1 Gusong-dong, Yusong-gu, Daejeon, 305-701, S. Korea  aemail : sjlee08@kaist.ac.kr, bemail : dkkim@kaist.ac.kr  Keywords: UHTCs, Hot-pressing, ZrB2, SiC, Oxidation behavior  Abstract. Zirconium diboride based composites containing silicon carbide with relative densities in excess of 99 % were produced by hot-pressing. Oxidation test was conducted in air at 1500 °C. ZrB2-SiC composite showed relatively low oxidation resistance due to the non-uniform surface silica-rich layer. But in case of mixed boride-SiC composites further improvement of the oxidation performance were observed due to the phase separation in the surface silica-rich layer.  Introduction Ceramic borides, carbides and nitrides are characterized by high melting points, high mechanical properties and relatively good oxidation resistance in extreme environments. This family of ceramic materials is known to be as Ultra High Temperature Ceramics (UHTCs) [1, 2]. The oxidation of ZrB2 has been studied by number of investigators [3, 4]. Many attempts have been made to enhance the oxidation resistance of ZrB2-based materials through appropriate additives. Considerably, the most common additive is SiC, which enhance oxidation resistance by forming a silica-rich scale [4, 5] and limits diboride grain growth [6]. Modification of the chemical composition of the oxide surface layer is one of the efficient ways of controlling oxidation resistance of non-oxide ceramics [2]. It is known from the literature that borate and silicate glass containing group IV-VI transition metal oxide show strong tendency to phase separation [7]. Therefore, this study has been focused on the improvement of the oxidation resistance by modifying surface glass scale.   Experimental Procedure  Processing. Commercially available ZrB2 (Grade A, H.C Starck), NbB2 (Japan New Metal Co., LDT), TaB2 (Japan New Metal Co., LDT) and SiC (H.C. Starck, UF25) were used to prepare the materials. Three compositions, their designation, and processing history are summarized in Table 1. To reduce the particle size in some batches, as-received ZrB2, NbB2 and TaB2 were vibration milled, respectively. Powders were milled in ethanol for 2 h using steel balls. During the milling 4 wt. % of Fe impurity was introduced because of wear of the steel balls. Fe impurity was successfully removed by an acid treatment. Subsequently, milled powders were ball-milled in ethanol for 24 hours with SiC. The solvent of the powder mixtures were dried on hot-plate with continuous strring, and the resulting dried powder mixtures were ground. Thus, obtained milled powders were hot-pressed (Thermal Technology Inc, Astro Hot Press) in BN-coated graphite dies. Powder compacted were heated under Ar flow at 1800 °C with an average heating rate of ~20°C/min. When the die temperature reached 1800 °C, a uniaxial load of 32 MPa was applied and dwelled for 2 h at 1800 °C, and then the furnace was cooled at a cooling rate of ~20 °C/min to room temperature. Oxidation. The oxidation test was carried by exposing the sample under air atmosphere after diamond-polishing to a 1 Dm using routine metallographic methods. Prior to oxidation, specimens were cleaned with acetone in an ultrasonic bath. Each specimen was heated a heating rate ~5 °C/min to the 1500 °C and held for 30 min in a tube furnace.  \\x0c', '254  SiAlONs and Non-oxides  Table. 1. Summary of UHTC compositions, designations, processing and relative density Composition Designation Processing Density (%)a ZrB2 - 30 vol. % SiC ZS 1800°C, 32 MPa 99  (Zr0.7Nb0.3)B2 - 30 vol. % SiC ZNS 1800°C, 32 MPa 99 (Zr0.7Ta0.3)B2 - 30 vol. % SiC ZTS 1800°C, 32 MPa 99 a Based on rule-of-mixture   Result and discussion  Bulk densities for the hot pressed billets were measured and the results showed 5.27 g/cm3 (ZS), 5.3 g/cm3 (ZNS) and 6.29 g/cm3 (ZTS), respectively. Using a rule of mixture calculation, and assuming that the true densities 6.09 g/cm3 for ZrB2, 6.61 g/cm3 for NbB2, 11.7 g/cm3 for TaB2 and 3.21 g/cm3 for of SiC. Based on this true density, all samples had >99% relative densities. XRD showed (data not shown) metal boride - SiC composites were successfully fabricated at 1800 °C. In case Nb, Ta doping, the slight shift of peaks was observed, which clarify the solid solution of Nb and Ta atoms into the Zr lattice site. SEM analysis did not reveal any obvious porosity in the microstructure, which supports the results of the density measurement and vibration milling was effective to enhance sintering driving force due to the reduced particle size. In case of ZS, average grain size of ZrB2 was 2~6 µm while average grain size of ZNS, ZTS were 0.8~5µm and 0.6~4 µm respectively. This result indicated that by adding NbB2 and TaB2, average grain size of ZrB2-transition metal boride solid solution became small.  Oxidation of three samples at 1500 °C for 30 min in air produced structures that consist of three layers: (1) surface silica-rich layer, (2) a SiC-depleted layer, and (3) un-reacted layer. The formation of three layers is consistent with previous investigation [5, 8, 9]. Fig. 1 is a SEM image of the top surface of the oxidized materials at 1500 °C, for 30 min.  (a)  (b)  (c)  Fig. 1. SEM images of oxidized top surface of materials: (a) ZS, (b) ZNS, and (c) ZTS \\x0c', 'Key Engineering Materials Vol. 403  255  It is shown that in non-uniform surface silica-rich layer was observed. This may be due to wetting characteristics or low viscosity of surface silica-rich layer owing to relative low temperature compared with melting point of SiO2 that might enhance the local oxidation rate. But in case ZNS and ZTS, surface silica-rich layer was uniformly distributed and expected parabolic oxidation behavior. Materials containing NbB2, TaB2 might be less oxidation than ZS, which is the result of phase separation in surface silica-rich layer induced by Nb-, Tacontained oxide. Similar to previous investigation[10], the presence of light spots of surface silica rich layer in Fig. 1 (b) and (c) distributed in a silica-rich matrix is an indication of high temperature glass phase separation The light spots were Nb-, Tacontained oxide from the analysis by EDS. Both increased liquidus temperatures and viscosities, which are characteristic features of phase separations, are good for a decrease in oxygen diffusisity through surface silica-rich layer. And cross-sectional analysis showed (data not shown) similar results with top surface analysis. ZS showed relatively thick surface silica-rich layer and SiC-depleted layers due to the non-uniform surface Si containing layer. In contrast, ZNS and ZTS showed thin surface silica-rich layer and SiC-depleted layer due to the uniform surface silica-rich layer.    Summary  Billet of ZrB2 containing SiC particulate addition of 30 mol % of NbB2 and TaB2 were produced using hot pressing of commercial powders. Nearly full densification was achieved relatively low temperature due to the size reduction of starting borides powder via vibration milling. In case of ZS, relatively poor oxidation resistance was observed due to the non-uniform distribution of surface silica rich layer. But in case of ZNS and ZTS, improved oxidation resistance was observed due to the uniform distribution of surface silica-rich layer. Oxidation of the mixed boride borides (ZNS, ZTS) results in the formation of corresponding oxide in the surface silica-rich layer. It should be emphasized that the oxidation resistance of mixed boride was higher than that of that ZrB2 ceramics due the phase separation phenomena.  References [1] S. R. Levine, E. J. Opila, M. C. Halbig, J.D. Kiser, M. Singh, and J. A. Salem: J. Eur. Ceram. Soc., Vol. 22 (2002), P. 2757 [2] M. M. Opeka, I.G Talmy, and J. A. Zaykoski: J. Mater. Sci., Vol. 39 (2004), P. 5887 [3] A. K. Kuriakose and J. L. Margrave: J. Electrochem. Soc., Vol. 111 (1964), P. 827 [4] W. C. Tripp and H. C. Graham: J. Electrochem. Soc. Vol. 118 (1968), p.1195 [5] A. Rezaie, W. G. Farenholtz, and G. E. Hilmas: J. Am. Ceram. Soc. Vol. 89 (2006), P. 3240 [6] A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas: J. Am. Ceram. Soc. Vol. 87 (2004), p. 1170 [7] W. Vogel: Glass Chemistry (Springer-Verlag, NewTock, 1994) [8] A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas: J. Eur. Ceram. Soc. Vol. 27 (2007), P. 2495 [9] S. S. Hwang, A. L. Vasiliev, and N. P. Padture: Mat. Sci. Eng. A Vol. 464 (2007) P. 216 [10] I. G. Talmy, J. A. Zaykoski, M. M. Opeka, and A. H. Smith J. Mater. Res. Vo. 21 (2006), P. 2593 \\x0c', 'SiAlONs and Non-oxides   10.4028/www.scientific.net/KEM.403   The Oxidation Behavior of ZrB2-Based Mixed Boride   10.4028/www.scientific.net/KEM.403.253       \\x0c']"
},{
  "_id": 264,
  "PDF": "The thermal stability in air of hot-pressed diboride matrix composites for uses at ultra-high temperatures.pdf",
  "Text": "['Corrosion Science 47 (2005) 2020-2033  www.elsevier.com/locate/corsci  The thermal stability in air of hot-pressed  diboride matrix composites for uses at  ultra-high temperatures  Fre´de´ ric Monteverde *  Institute of Science and Technology for Ceramics, National Research Council,  Via Granarolo, 64-48018 Faenza, Italy  Received 19 March 2004; accepted 24 September 2004  Available online 7 December 2004  Abstract  The  resistance  to oxidation in ambient air at a temperature up to 1600 °C of  two hot pressed diborides matrix composites, both containing 19.5% v/o SiC and 3 v/o HfN (as sinter ing aid), was investigated. The diboride matrix was based on HfB2 or a ZrB2/HfB2 mixture (volume ratio \\x19 1). Both the materials were subjected to repeated heating-cooling cycles at 1600 °C, and a 20 h exposure at 1450 °C in flowing dry air. Modest weight gains and limited  good  thermal  stability.  In  accordance with  the  corrosion depths highlighted a rather thermo-gravimetric test at 1450 °C,  the oxidation kinetics  for both the composites  superbly  fit a para-linear law. The introduction of  the SiC particles provided tangible benefits for the  resistance  to oxidation. One of  the oxidation products, a borosilicate  glass,  sealed pores  and coated the exposed faces, greatly limiting the inward transport of oxygen towards  the  internal oxide/diboride interfaces.  Ó 2004 Elsevier Ltd. All rights reserved.  Keywords: UHTC; Hot-pressing; ZrB2; HfB2; Microstructure; Resistance to oxidation  * Tel.: +39 0546 699 758; fax: +39 0546 463 81.  E-mail address: fmonte@istec.cnr.it  0010-938X/$ see front matter Ó 2004 Elsevier Ltd. All rights reserved.  doi:10.1016/j.corsci.2004.09.019  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2021  1. Introduction  Ultra-high temperature ceramics (UHTCs) can be defined as materials that have uses at temperatures in excess of 1800 °C. Ceramics that meet this classification may  include diborides and carbides of early transition metals like hafnium zirconium or  tantalum [1]. In short, these materials are expected to behave superbly as thermal shieldings because of their unrivaled melting points exceeding 3000 °C, coupled to re tained strengths and a thermal stability at very high temperatures [2,3]. In the future,  the employment of UHTCs  is expected to have a broad impact on several  space  applications, ranging from thermal protection structures (TPSs)  for modern space vehicles  to propulsion for  rocket motors or  jet  engines  [4]. Concrete perspectives  are currently addressed towards the re-design of a new generation of re-usable space  vehicles with very sharp (not actively cooled) leading edges [5]. This innovative class  of materials has the potential to overturn the aerodynamic constraint that only TPSs  with blunted profiles can endure the peak temperatures generated when a vehicle through the EarthÕs atmospheric gases during hyper-sonic flights  repeatedly cuts  [6]. The highly-localized heating loads  foreseen in the  interaction regions dictate  the need for outstanding key structural  issues like oxidation and ablation resistance.  Even in the absence of ablation, dissociated atmospheric gases may diffuse to the  vehicle surfaces,  inducing additional heating rates associated with the atomic heat  of  formation. Therefore, a prominent  surface catalytic efficiency,  i.e.  the capacity  to conduct away or re-radiate excess energy,  is also strongly demanded [6].  Current research activities still encounter difficulties in manufacturing near net shape dense components. Pressure-assisted sintering at temperatures approaching 2000 °C are required to prevail over the intrinsically poor sinterability of  the dibo rides matrices  [7]. However,  such conditions are known to induce  coarsening of  the microstructure and a general degradation in the thermo-mechanical stability.  In this  sense,  the  scientific  community has been addressing the deficiencies of  UHTCs, for instance the poor sinterability and the resistance to oxidation/ablation,  with an appropriate focus on compositional designs and composite constructions  [4,8-12]  in order to corroborate more credibly the applicability of this class of TPSs  to re-usable future space-crafts.  The present paper focused on the response to oxidation of two diborides matrix  composites containing 19.5 v/o SiC and 3 v/o HfN (as sintering aid). Their thermal  stability in ambient air was studied monitoring the weight change in accordance with  specific heat treatments. The modifications in the microstructure were examined and  correlated to the oxidation mechanisms involved.  2. Experimental  2.1. The fabrication of  the composites  The starting powder mixtures (v/o), HfB2 + 19.5 b-SiC + 3 HfN (material HS) and ZrB2 + 37.5 HfB2 + 19.5 b-SiC + 3 HfN (material ZHS) were milled for one  \\x0c', '2022  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  day in a polyethylene jar using absolute ethanol and zirconia balls, and then dried  and sieved. The as-processed powder mixes were hot-pressed at low vacuum (0.5 mbar), 8 °C/min heating rate, with an applied pressure of 50 MPa using a BNthe hot-press run was 1900 °C  lined induction-heated graphite die. The set point of  for both the compositions, whilst  the soaking time was 40 and 30 min for material  HS and ZHS,  respectively. Some thermo-mechanical properties of  the composites  are listed in Table 1. Densification behavior, microstructure, the role of the additive  and thermo-mechanical properties are commented on in detail elsewhere [13].  2.2. The oxidation treatments  The  composites were  subjected  to  various  oxidation  treatments  (at  ambient  pressure):  (T-1)  (T-2)  (T-3)  non-isothermal run up to 1450 °C in flowing dry air (50 cm3/min), heating rate 2 °C/min and free cooling; isothermal run at 1450 °C for 20 h in flowing dry air (50 cm3/min), heating rate 30 °C/min and free cooling; isothermal runs at 1600 °C, cumulating in 20, 40 and 60 min of exposure (load ing  and removal of  the  coupon at  the fixed set-point  for  each scheduled  exposure).  Coupons with dimensions of 2.5 · 2.0 · 10.0 mm3 were cut out from the sintered (surface finish Ra 0.2 lm), washed in an ultrasonic bath of acetone, and then billet dried at 80 °C overnight. Treatments T-1 and T-2 were executed using a thermo gravimetric analyzer \\x003 mg of accuracy,  10  (mod.  STA409, NETZSCH Gera¨ tebau GmbH-Germany),  equipped with a vertically-heated Al2O3  chamber. The  test  pieces were placed on zirconia supports,  separating them from the Al2O3 holder.  In treatment T-2, the application of a fast heating-rate prior to the isothermal period the temperature of 1450 °C was set.  aimed to minimize oxidation effects until  Treatment T-3 was performed using a bottom-loading furnace box, heated with  MoSi2  elements, and insulated with highly-porous Al2O3 fiber. The  test  coupons  were placed upon a reaction-bonded sintered SiC support, ensuring the minimal con tact area between them. The sample mass was measured before and after each cycle.  2.3. The analysis of  the microstructure  The microstructure was analyzed using a scanning electron microscope (SEM,  mod. S360, Cambridge, UK)  equipped with an energy dispersive microanalyzer  (EDX, mod. INCA Energy 300, Oxford Instruments, UK), and an X-ray diffracto meter (XRD, mod. D500, Siemens, Germany). Sections of some oxidized coupons (finished 0.25 lm), and then observed  were first polished using diamond abrasives  via SEM-EDX analyses. These cross-sections were imaged using secondary electrons  (SEs),  free of conductive coating in order  to maintain the sensitivity of  the EDX  equipment  to the low Z elements as high as possible.  \\x0c', 'Table 1  Thermo-mechanical properties of the materials HSH and ZHS  Composite  aE (GPa)  bHv1.0 (GPa)  p  ffiffiffiffi  bKIC (MPa  m  )  HS  ZHS  506 ± 4  497 ± 4  22.3 ± 0.9  22.0 ± 0.7  CN  3.75 ± 0.30  -  DCM  6.2 ± 0.5  5.1 ± 0.8  \\x006/°C) k (10  25-800 °C  7.16  6.45  25-1000 °C  25-1300 °C  7.25  6.70  7.21  6.77  br (MPa)  25 °C  650 ± 55  765 ± 20  1500 °C  465 ± 45  250 ± 45  Youngs modulus E, micro-hardness Hv1.0, fracture toughness KIC (CV: chevron notch; DCM: direct crack measure), flexural strength r. a Uncertainity. b Mean ± 1std. dev.  linear thermal expansion coefficient k,  F  .  M  o n  t  e v e  r  d  e  /  C  o  r r  o  s  i  o n  S  c  i  e  n  c e  4 7  ( 2 0 0 5 )  2 0 2 0 - 2 0 3 3  2 0 2 3  \\x0c', '2024  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  3. Results  3.1. The microstructure of  the hot-pressed composites  On the basis of the SEM observations on typical polished sections, the final rela tive density of  the hot-pressed composites was deemed greater than 99.5% (Figs. 1  and 2). Moreover,  the general aspects of  the microstructure of both the materials (from 0.5 to 4 lm of grain  are  characterized by regularly-faceted diboride grains  size), and several original contact points among diboride grains and the SiC partic ulate. The latter phase, the majority of darker features in Figs. 1 and 2, is distributed  intergranularly. Relevant defects were not revealed. In material HS,  the XRD ana lysis detected, apart HfB2 and SiC, a cubic Hf(C,N) solid solution as secondary crys talline phase. Conversely, the XRD outcome from material ZHS provided evidence,  SiC apart, of diverse hexagonal  (Zr,Hf)B2  solid solutions  and a minority  cubic  (Zr,Hf)(C,N) solid solution. The crystallochemical specifics of  the hexagonal HfB2  and ZrB2 were expected to allow readily mutual solubility.  Fig. 1. SEs-SEM micrograph from a polished cross-section of the composite HS.  Fig. 2. SEs-SEM micrograph from a polished cross-section of the composite ZHS.  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2025  The addition of HfN as sintering aid favored densification of these highly covalent  matrices and enabled full density being achieved. The SiC particles hindered exces sive growth of the diboride matrix and kept the average grain size at less than 2 lm. In both the as-sintered composites, other reaction phases of the hot-pressing,  secondary products of minor  importance, were  identified. A detailed discussion  regarding the microstructure development in the composite HS and ZHS is reported  elsewhere [13].  3.2. The response to oxidation  The specific mass change (w) vs. temperature (T) during the treatment T-1 is depicted in Fig. 3. An heating rate of 2 °C/min enabled measuring peculiar tempera tures,  for  instance when the  sample began gaining weight appreciably (Table 2).  The wavy pattern of  the curves is closely connected to a certain thermal  instability  of the oxide scale covering the external faces of the oxidizing sample. The decreasing  branch of the thermo-gravimetric (TG) data between Tmin and Tmax was already de scribed on similar diborides-SiC systems [14,15]. The slowing down of  the weight gain data can be ascribed to the  release of B2O3  (g) as a temporary dominating  mechanism of  the oxidation process.  Fig. 3. Weight change (w) vs.  temperature, recorded during the treatment T-1. Characteristic values of  temperatures Ton, Tmax and Tmin are listed in Table 2.  Table 2  Thermal treatment T-1: characteristic temperatures (see Fig. 3) and weight change (w)  Composite  HS  ZHS  Ton (°C)  735  700  Tmax (°C)  1140  1130  Tmin (°C)  1300  1260  w (mg/cm2)  0.55 ± 0.02  0.80 ± 0.02  \\x0c', '2026  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  Fig. 4. Thermo-gravimetric data (w) vs. exposure time (t), HS. The multiple linear regression fit (A) is shown.  in accordance to the treatment T-2 of material  Fig. 5. Thermo-gravimetric data (w) vs. exposure time (t), ZHS. The multiple linear regression fit (A)  is shown.  in accordance to the treatment T-2 of material  In accordance with the treatment T-2, isothermal oxidation tests were further accomplished (Figs. 4 and 5). An offset of 0.25 mg/cm2, which accounts for the oxi dation preceding the planned permanence, was equally subtracted from both the TG  curves. Supported by the multiple-law model proposed by Nickel [16], some interestp þ K LIN t, of the weight change data (w) along the exposure time (t) law, w ¼ K PAR ing results emerged. In both cases, the multiple linear-regression fit over a para-linear yielded excellent solutions (R2 > 99.9). The most relevant findings of this calculation  t  ffiffi  were that  the oxidation process is rate-limited by diffusional mechanisms and that  the positive linear contribution is not negligible (Table 3). An extra logarithmic con tribution,  KLOG log(t),  was  discharged  because  a  (calculated)  negative  KLOG  throughout  the 20 h of exposure is deprived of physical meaning.  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2027  Table 3  Treatment T-2: weight gain (w) and calculated constants of  the multiple linear regression fit  (see Figs. 4  and 5)  Composite  HS  ZHS  w (mg/cm2)  2.80  4.35  KPAR  \\x004 h \\x001) (mg2 cm  a0.2  0.635  KLIN  (mg cm  \\x002 h \\x001)  0.0338  0.0385  \\x004 h \\x001 [17]. a For comparison: KPAR = 0.8 mg2 cm  Finally, the thermal treatment T-3 was carried out: after the loading of the samtemperature of 1600 °C. The  the furnace took about 15 min to reach the exact  ple,  cumulative mass  changes  are  reported  in Table  4. After  the  first  exposure  of  20 min, a linearly increasing trend of  the weight gain seems to prevail.  3.3. Microstructural changes after oxidation treatments  As far as treatment T-3 is concerned, the micrographs in Fig. 6 illustrate the modification of the virgin microstructure after the repeated exposures at 1600 °C. On top  Table 4  Thermal treatment T-3: (cumulative) mass change (w), 20 min of exposure each  Composite  HS  ZHS  w (mg/cm2)  20 min  1.30 ± 0.05  3.40 ± 0.05  (2 · 20) min  2.00 ± 0.05  4.60 ± 0.05  (3 · 20) min  2.90 ± 0.05  5.30 ± 0.05  Fig. 6. SEs-SEM micrographs from polished cross-section of material HS (a) and ZHS 3 · 20 min of exposure at 1600 °C (treatment T-3). Arrows mark the outermost glass layer.  (b),  after  \\x0c', '2028  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  of the exposed faces, an evident glassy scale has formed: it basically consists of silica.  The undulating trend of the thickness of such glassy product indicates that, owing to  a diminished viscosity at  the testing temperature,  it may laterally flow out. Under neath, an additional oxide  scale which separates  the outermost glassy scale  from  the as-sintered virgin bulk is basically constituted of monoclinic HfO2 or (Zr,Hf)O2  in the composite HS and ZHS, respectively. These oxide crystals are enclosed by a  glassy melt. The patchwork of micrographs in Fig. 7 shows a representative cross section of the oxidized composite HS. The diboride-SiC skeleton of the material tol erates the applied thermal  load well. A very similar configuration can be drawn for  the composite ZHS, apart from an enhanced penetration of the oxidation front. ItÕs worth noting that a secondary phase like Hf(C,N) exhibited a certain thermal  instability:  the signature of an incoming oxidation is rather evident  (see the circles  in Fig. 7).  The SEM-EDX investigation of 1450 °C (treatment T-2)  at  the cross-sectioned samples tested isothermally  confirmed the basic  assemblage previously described  for treatment T-3 (Fig. 8). Obviously, the longer exposure during treatment T-2 fa vored the oxygen penetrating more deeply into the bulk,  i.e. more-widely corroded  Fig. 7. Patchwork of SEs-SEM micrographs from a polished cross-section of material HS, after thermal treatment T-3 (3 · 20 min).  In the upper part: bottom of  the oxide scale. Circles  locate partly oxidized  Hf(C,N) grains.  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2029  Fig. 8. SEs-SEM micrographs from polished cross-section of material HS (a) and ZHS (b), after thermal  treatment T-2.  Fig. 9. XRD pattern from an oxidized surface of material ZHS, after  thermal  treatment T-2. Lattice  parameters a, b, c, and b extrapolated using a Rietveld refinement) and Miller indexes are shown.  internal  regions. The undulating variation of  the glass thickness  induced a similar  pattern of  the corroded portions in both the composites. The XRD analysis of  the  exposed surfaces detected monoclinic HfO2 or a (Zr,Hf)O2 solid solution in the com posite HS and ZHS, respectively. In addition,  the characteristic hump in the XRD  pattern confirmed the existence of a surface glassy product. Only for  the oxidized  material ZHS,  the  lattice  parameters  of  a  (monoclinic)  structure,  space  group  P2_1/c, were estimated using a Rietveld refinement [18]: as expected, they ranged in side the interval of  the known cell parameters of  the boundary phases ZrO2 and  HfO2  (Fig.  9). Local  ruptures  of  the  external  oxide  scale were  not  observed.  The EDX analysis of  the  glass basically  assessed a  chemical  composition of  a  \\x0c', '2030  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  borosilicate. Elemental Hf, Zr or C were not detected. Surprisingly, a significant  fraction of the SiC particles lying close to the outer oxide layer appeared only slightly  affected by the oxidation attack (Fig. 7).  4. Discussion  First of all, the thermo-gravimetric output of oxidation processes involving con current gain and loss of mass (which is our case) was interpreted only in close agree ment with the microstructural evidence relative to the oxidized samples. The present  study in fact highlighted that rather limited mass gains were accompanied by a few  alterations  to the original microstructure. On the other hand,  the response of  the  refractory materials  to high-temperature oxidizing conditions  imposed by the cur rently-used furnace heating is known to differ considerably from those experienced  during the Earth descent of a space mission. As regards the prospects of the appli cability of a re-entry-simulating apparatus,  for  instance an arc plasma jet  reactor,  the oxidation resistance of  the composite was anyhow investigated in accordance  with well-established conventional procedures.  The results reported in the previous section showed that, within the applied test ing conditions, the response to oxidation in laboratory air of both the diborides-SiC  composites was rather good. Specifically, apart  from the intrinsic refractoriness of  the diboride matrices, the introduction of about 20 v/o SiC particles actively contrib uted to an increase in the resistance to oxidation. Moreover, the adoption of an HfB2  matrix instead of a ZrB2/HfB2 mixture, whatever the experimental conditions, appre ciably enhanced the overall oxidation resistance (Tables 2-4).  Depending  on  the  temperature  range  selected,  the  expected main  reactions  describing the oxidation process, M = Hf or (Zr,Hf), MB2 þ 5=2O2 ¼ MO2 þ B2O3 ðlÞ  SiC þ 3=2O2 ¼ SiO2 þ COðgÞ  MC1\\x00xNx þ ð3 \\x00 xÞ=2O2 ¼ MO2 þ ð1 \\x00 xÞCOðgÞ þ x=2N2 ðgÞ  B2O3 ðlÞ ¼ B2O3 ðgÞ  SiO2 ðsÞ þ COðgÞ ¼ SiOðgÞ þ CO2 ðgÞ  2SiO2 ðsÞ þ SiCðgÞ ¼ 2SiOðgÞ þ COðgÞ  ð1Þ  ð2Þ  ð3Þ  ð4Þ  ð5Þ  ð6Þ  may be all, or some of them, simultaneously active. The reactions (1)-(3) and (4)-(6)  involve mass gain or mass loss, respectively. Other equilibrium equations, candidate  mechanisms for the formation/decomposition of silica were omitted for the sake of  brevity [19]. The main reaction products from the oxidation of MB2, which consti tutes the skeleton of the tested composites, are MO2 and amorphous B2O3 (reaction latter phase has a very low melting point, 450 °C, and an high vapour  (1)). The  pressure:  these  two conditions  favor  its volatilisation (i.e.  loss of mass) at high  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2031  temperature. The increase in the scaling rate above 700 °C (Fig. 3) was primarily  caused  by  the  selective  oxidation  either  of  the MB2  (reaction  (1))  and  of  the  M(C, N) secondary phase (reaction (3)), with no appreciable attack of the SiC par ticles. The M(C,N) grains directly facing or near the ambient atmosphere fairly oxi dized and led an incipient porosity to form. The oxide scale that  forms partially  allows the diffusion of oxygen through interconnected pores or via lattice vacancies  of MO2. Similarly to ZrO2 or HfO2, (Zr,Hf)O2 behaves plausibly as an anionic con ductor. In addition, at relatively high temperature the (fluid) B2O3, which most likely  covers the external faces of the oxidizing sample,  is known to be much more perme able to oxygen than the silica glass [20]. The decreasing trend of the TG data between  Tmin and Tmax (Fig. 3) primarily accounted for the noticeable loss of mass through  the volatilisation of B2O3(g).  Substantial benefits from the presence of the SiC particles resulted only after increasing the temperature above 1400 °C. It is believed that a number of SiC parti cles were present at or near the ambient MO2 interface following the transitory heat ing stage. In accordance with reaction 2, silica which forms from the SiC particles  spreads laterally and across the surface of the MO2 grains, combines with the avail able boria and provides a borosilicate layer more protective than the MO2 alone.  Likewise other diborides/SiC systems [4,11,14,15], the exposed faces of the coupons  were coherently covered by an adherent borosilicate glass, even though its undulat ing pattern was somehow responsible for the deeper infiltration of oxidation front in  correspondence to the thinner section of  the protective glassy layer (Fig. 8).  As far as the isothermal treatment T-2 is concerned, available thermo-gravimetric  referenced data for comparison are really scarce. A well-researched campaign of TG  oxidation tests over an (88% dense) HfB2-SiC composite was attempted [17]. The  constitution of the multilayered corroded structures depicted in Fig. 8 basically agreed with the evidence of Hinze et al. which, whatÕs more,  indicated the diffusion  of oxygen through the silica glass as the rate-limiting step of a pure parabolic oxida tion (in flowing pure oxygen). In our case, both the sets of TG data superbly fits a  para-linear law (Figs. 4 and 5). The preponderant parabolic contribution dominates  the monotonically decelerating trend of the TG data. In fact, the development of a  protective external oxide scale progressively imposes longer diffusion paths for oxy gen to arrive at  the diboride-oxide interface. The departure from a pure parabolic  pattern is described by a little linear contribution. Such additional linear dependence  of the weight change vs. the oxidation time often characterized the response to oxi dation of multiphase non-oxide ceramics. The motivations for this deviation from a  plain parabolic model may depend on the simultaneous progress of several physical  factors, for instance the reaction—diffusion-vaporization of gaseous oxidation prod ucts. The ultimate assessment of such deviation is still a matter for study.  The presence of SiC as particulate was certainly responsible for the rather good oxidation resistance. At 1450 °C,  intergranular SiC particles  the majority of  the  began oxidizing, yielding silica as an oxidation product, and continuously contribut ing to form a borosilicate glass. This melt, scarcely viscous at the testing temperature  and characterized by a restricted permeability to oxygen, easily sealed the external  surfaces  of  the  test  piece  and  short-circuited  paths  for  the  incoming  oxygen  \\x0c', '2032  F. Monteverde / Corrosion Science 47 (2005) 2020-2033  (i.e. residual porosity, cracks) towards accessible MB2 grains and/or grain boundary  phases prone to rapid oxidation. The evidence of residual un-oxidized SiC particles  just beneath the external oxide scale (Fig. 7) settled the fundamental role of such a  phase in suppressing the advance of  the oxidation attack.  In contrast  to previous  studies on similar systems [4,15,17,21], a partial depletion of SiC at high temperature  and reduced oxygen partial pressure in consequence of an active oxidation mecha nism was not appreciated. Moreover,  the apparent cleanness of  the diboride-dibo ride  boundaries  (i.e.  absence  of  intergranular  compounds)  acted  as  a  positive  feature in slowing down the preferential  inward transport of oxygen through them.  Furthermore,  the SiC particulate, mostly residing at  triple points of  the diboride  grains, additionally contributed to block the inward access of the oxygen, and there fore  the  conversion of MB2  into MO2. However,  the  secondary grain boundary  phase like M(C,N) shows it  to be a weak point against oxidation: under oxidizing  conditions in fact it tends to degrade rapidly, and to introduce preferential paths for further provision of oxygen into the materialÕs inner regions.  5. Summary  Two highly-dense diboride matrix composites, containing 19.5 v/o SiC and 3 v/o  HfN as sintering aid, were successfully hot-pressed. The diboride matrices consisted  of HfB2 or a mixture of ZrB2 and HfB2. The resistance to oxidation was tested in  laboratory air, monitoring not only the weight change but also the modifications  in the microstructure induced by the oxidation. For both the composites, the thermotest at 1450 °C for 20 h obeyed para-linear  gravimetric data along an isothermal  kinetics. The parabolic contribution dominates the monotonically decelerating pat tern of the thermo-gravimetric data. The main oxidation products were monoclinic  MO2, M = Hf or (Zr,Hf), a borosilicate glass, and volatile B2O3. The release of gas eous oxidation products had an appreciable impact for the thermal treatments below 1350 °C. Above 1400 °C,  the SiC particles markedly improved the  the presence of  resistance to oxidation of both the composites. A protective borosilicate glass layer,  scarcely viscous at the testing temperatures, sealed pores and efficaciously coated the  sample surfaces. The intergranular distribution of the SiC particulate proved to be a  key feature in pinning the oxidation front advancing along the diboride/diboride  grain boundaries apparently free of secondary phases. Given a limited susceptibility (up to 1600 °C) to experience some small structural damage, extra studies should be  further addressed towards the validation of these UHTCs under Earth re-entry mis sion conditions.  Acknowledgement  The author acknowledges the very valuable support of Dr. A. Balbo for the ther mo-gravimetric tests.  \\x0c', 'F. Monteverde / Corrosion Science 47 (2005) 2020-2033  2033  References  [1] A. Bronson, Y.-T. Ma, R.R. Mutso, N. Pingitore, Compatibility of  refractory metal boride/oxide  composites at ultra high temperatures, J. Electrochem. Soc. 139 (11) (1992) 3183-3196.  [2] K. Upadhya,  J.-M. Yang, W.P. Hoffmann, Materials  for  ultra-high  temperatures  structural  applications, Am. Ceram. Soc. Bull. 58 (1997) 51-56.  [3] C.R. Wang, J.-M. Yang, W.P. Hoffmann, Thermal stability of refractory carbide/boride composites,  Mater. Chem. Phys. 74 (2002) 272-281.  [4] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra high  temperature ceramics for aeropropulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757-2767.  [5] Slender Hypervelocity Aerothermodynamic Research Probes (NASA Project).  [6] P.A. Gnoffo, Planetary entry gas dynamics, Annu. Rev. Fluid. Mech. 31 (1999) 459-494.  [7] H. Pastor, Metallic borides: preparation of solid bodies—sintering methods and properties of solid  bodies,  in: V.I. Matkovich (Ed.), Boron and Refractory Borides, Springer Verlag, New York, 1977,  pp. 457-493.  [8] J.D. Bull et al.,  in: US Patent 5,750,450 Ablation resistant Zirconium and Hafnium Ceramics, 1998.  [9] M. Opeka,  I.G. Talmy, E.J. Wuchina,  J.A. Zaykoski,  S.J. Causey, Mechanical,  thermal  and  oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. Soc. 19 (1999)  2405-2414.  [10] F. Monteverde, A. Bellosi, Advances  in microstructure  and mechanical properties of  zirconium  diboride-based ceramics, Mat. Sci. Eng. A 346 (2003) 310-319.  [11] A. Bellosi, F. Monteverde, Fabrication and properties of zirconium diboride based ceramics for UHT  applications, in: Proceedings 4th European Workshop on Hot Structures and TPS for Space Vehicles,  ESA-SP521, April 2003, pp. 65-71.  [12] F. Monteverde, A. Bellosi, Microstructure and properties of an HfB2-SiC composite for ultrahigh  temperature applications, Adv. Eng. Mater. 6 (5) (2004) 331-336.  [13] F. Monteverde, A. Bellosi, The  efficacy of HfN as  sintering aid in the manufacture of ultrahigh  temperature diborides matrix ceramics, J. Mater. Res. 19 (12) (2004) 3576-3585.  [14] F. Monteverde, A. Bellosi, Oxidation of ZrB2 based ceramics in dry air, J. Electrochem. Soc. 150 (11)  (2003) B552-B559.  [15] F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC composite, J. Eur. Ceram.  Soc.,  in press.  [16] K.G. Nickel, Multiple law modelling for the addition of advanced ceramics and a model-independent  figure  of merit,  in: K.G. Nickel  (Ed.), Corrosion  of Advanced Ceramics/Measurements  and  Modelling, Kluwer Academic Publishers, Dordrecht, The Netherlands, 1994, pp. 59-71.  [17] J.W. Hinze, W.C. Tripp, H.C. Graham, The high temperature oxidation behaviour of a HfB2 +  20 v/o SiC composite, J. Electrochem. Soc. 122 (9)  (1975) 1249-1254.  [18] PowderCell for Windows 2.4, W. Kraus and G. Nolze (BAM-Berlin, Germany).  [19] G. Hilfer, The impact of environmental conditions during re-entry of  the re-usability of a Si-based  ceramic TPS, in: Proceedings 4th European Workshop on Hot structures and TPS for Space Vehicles,  ESA-SP521, April 2003, pp. 363-368.  [20] R.H. Doremus, Glass Science, John Wiley and Sons, New York, 1973, p. 138.  [21] W.C. Tripp, H.H. Davis, H.C. Graham, Effect of an SiC addition on the oxidation of ZrB2, Am.  Ceram. Soc. Bull. 52 (8) (1973) 612-616.  \\x0c']"
},{
  "_id": 265,
  "PDF": "The ZrB2 Volatility Diagram.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 88 [12] 3509 - 3512 (2005)  DOI: 10.1111/j.1551-2916.2005.00599.x  r 2005 The American Ceramic Society  The ZrB2 Volatility Diagram  William G. Fahrenholtz*,w  Materials Science and Engineering, University of Missouri-Rolla, Rolla, Missouri 65409  A volatility diagram was calculated for temperatures of 1000,  1800, and 2500 K to understand the oxidation of ZrB2. Applying the diagram, it can be seen that exposure of ZrB2 to air produces ZrO2 (cr) and B2O3 (l) over the temperature range considered. The pressure of the predominant vapor species was predicted to increase from B10 \\x006 Pa at 1000 K, to 344 Pa at 1800 K, and to B105 Pa at 2500 K. Predictions were consistent with experimental observations that ZrB2 exhibits passive oxidation below 1200 K, but undergoes active oxidation at higher  temperatures due to B2O3 (l) evaporation.  I.  Introduction  THE borides, metals are (UHTCs) because  carbides, and nitrides of  the  early transition  considered  ultra-high  temperature  ceramics  of melting  temperatures above 3000 K, attack.1,2 Among  high hardness,  and resistance  to chemical  the UHTCs, zirconium diboride (ZrB2) is a candidate for thermal protection systems and scramjet engine components for  hypersonic flight vehicles as well as high temperature electrodes, molten metal containment systems, and incinerators.3-6 Heating ZrB2 in air produces a scale composed of ZrO2 and B2O3.7,8 Below 1200 K, liquid B2O3 forms a continuous layer that wets the ZrO2 and the underlying ZrB2. The B2O3 (l) layer acts as a barrier to oxygen diffusion resulting in passive oxidation of t1/2) oxidation kinetZrB2 and parabolic (diffusion-limited or ics.9-11 At intermediate temperatures (1200-1700 K), the rates of  z  formation and volatilization of B2O3 (l) are similar, resulting in para-linear kinetics because of competition between mass gain  (ZrO2 and B2O3 formation) tion).12,13 Above 1700 K,  and mass  loss  (B2O3 active oxidation with rapid linear  vaporiza kinetics has been attributed to loss of B2O3 (l) by evaporation, which leaves behind a porous, non-protective ZrO2 (cr) layer.7,14 Interactions between gases and condensed phases can be interpreted with volatility diagrams.15 Volatility diagrams plot the  vapor pressure of the predominant gaseous species as a function  of oxygen partial pressure and temperature. Gas-solid interac tions have been studied for systems such as Mg-O, Al-O, Si-O, Si-C-O, Si-N-O, and Mg-O-C using volatility diagrams.16-18  A recent  study of UHTCs  employed volatility diagrams  for  metals (e.g., Zr, Si, B) to evaluate the thermal stability of oxide scales.19 However, to date, a true volatility diagram for ZrB2 has not been reported. The diagram is needed to understand the  oxidation of pure ZrB2 as well as the oxidation of ZrB2-based ceramics containing additives such as SiC, MoSi2, or graphite. The purpose of this paper is to describe the construction and  interpretation of a volatility diagram for ZrB2.  II.  Calculations and Diagram Construction  Several  thermodynamic databases were  examined to identify  relevant species containing Zr, B, and/or O, but only data from the NIST-JANAF tables were used to maximize consistency.20  After eliminating ionized species, duplicate data, and condensed  species that were not observed in oxidized specimens, 13 species  of interest were identified (Table I). Based on oxidation studies  reviewed in the introduction, the oxidation of ZrB2 (cr) to ZrO2 (cr) and B2O3 (l) by Eq. (1) was used to determine the equilibrium partial pressure of oxygen (pO2) for oxidation of ZrB2.  ZrB2  ðcrÞ þ 5  2O2  ðgÞ ! ZrO2  ðcrÞ þ B2O3  ðlÞ  (1)  Tabulated data were used to calculate the change in Gibbs’ free energy (DG0 rxn ) for Reaction (1) and for reactions that produced volatile species from ZrB2 (cr) or from ZrO2 (cr) and B2O3 (l). The DG0 rxn values were converted to equilibrium constant (Keq) values using Eq. (2), and then to equilibrium partial pressures  using expressions for the equilibrium constant for each reaction  such as the one presented as Eq. (3) for Reaction (1). Unit ac tivity was assumed for all condensed phases. The results are re ported as partial pressures (e.g., no units, assuming an ambient pressure of 1.013 \\x02 105 Pa or 1 atm). At 1800 K, the pO2 calwas 4.2 \\x02 10\\x0016 (vertical culated for the co-existence of ZrB2 (cr), ZrO2 (cr), and B2O3 (l) line in Fig. 1).  rxn ¼ \\x00RT ln Keq DG0  (2)  where R is the ideal gas constant and T is the absolute temper ature  K eq ¼  aB2 O3  ð ð  Þ aZrO2  ð ð  Þ  aZrB2  Þ aO2  Þ5=2 ¼  1  pO2  ð  Þ5=2  (3)  where a is the activity of the species involved in the reaction  Below the equilibrium pO2 for Eq. (1), the gases are in equilibrium with ZrB2 (cr). For example, B2O3 (g) can form by Eq. (4). Other gases that form by reaction of ZrB2 are listed in Table II.  ZrB2  ðcrÞ þ 5  2O2 ðgÞ ! ZrO2  ðcrÞ þ B2O3  ðgÞ  (4)  As pO2 increases, the amount of B2O3 (g) should increase since O2 is a reactant. This relationship can be seen in the portion of  3509  Table I.  Zr, B, and O Species of Interest for Calculation of the  Volatility Diagram  Zr species  Zr-O species  Zr-B species  B species  B-O species  Zr (g)  ZrO (g)  ZrB2 (cr)  B (g)  B2O3 (g) B2O3 (l) BO (g)  ZrO2 (cr) ZrO2 (g)  B2 (g)  BO2 (g) B2O (g) B2O2 (g)  A. Heuer—contributing editor  This material is based upon work supported by the National Science Foundation under  Grant DMR 034680.  *Member, American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: billf@umr.edu  z  Note: The NIST-JANAF convention is used whereby physical state is indicated in pa renthesis with (cr) for crystalline solids, (l) for liquids or amorphous solids, and (g) for gases.  Manuscript No. 20394. Received April 8, 2005; approved May 15, 2005.  \\x0c', '3510  Communications of the American Ceramic Society  Vol. 88, No. 12  ZrB2 (cr)  ZrO2 (cr) and B2O3 (l)  undergoes four major vapor transitions at 1800 K: (1) from B (g) to BO (g) at a pO2 of B2 \\x02 10\\x0027; (2) from BO (g) to B2O2 (g) at a pO2 of B3 \\x02 10\\x0017; (3) from B2O2 (g) to B2O3 (g) at a pO2 of B3 \\x02 10 \\x0016; and (4) from B2O3 (g) to BO2 (g) at a pO2 of B8 \\x02 101. Vapor pressure calculations were also completed at 1000 and 2500 K to show how the diagram changed with tem perature (Fig. 2(b)). Interestingly, the transition from stability of  B2O2 (g) to B2O3 (g) falls near the equilibrium pO2 for Eq. (1) (B4 \\x02 10\\x0016) at 1800 K and also at 2500 K so that the vapor transition is obscured by the transition in stable  condensed  phases.  BO2  B2O3  BO  B2O2  B2  B2O  B  Fig. 1.  Partial pressure of B species as a function of oxygen partial  pressure at 1800 K.  Fig. 1 to the left of the ZrB2/ZrO2-B2O3 line. Above the equilibrium pO2 for Eq. (1), the gases are in equilibrium with ZrO2 (cr) and B2O3 (l). In this regime, B2O3 (g) forms by direct vaporization of B2O3 (l) according to Eq. (5). ðlÞ ! B2O3  B2O3  ðgÞ  (5)  III.  Discussion  The volatility diagram was used to interpret experimental ob servations of ZrB2 oxidation in air. At 1000 K, the predominant species above B2O3 (l) and ZrO2 (cr) was BO2 (g) with a partial pressure of B10\\x0011 (B10\\x006 Pa). The next highest vapor pressure was B10 \\x0012 (B10 \\x007 Pa) for B2O3 (g). The calculated pressures of the other species were all dramatically lower (Table IV).  The vaporization rate for B2O3 (l) in air is expected to be low at 1000 K based on the partial pressures of the various gases, none \\x0011. of which exceeds B10 Above B1650 K, rapid, linear (active) oxidation kinetics have  been attributed to a significant increase in the rate of B2O3 (l) evaporation.7 Examination of ZrB2 oxidized at 1773 K by scanning electron microscopy (Fig. 3) and X-ray diffraction (not  shown) revealed that  the oxide layer was made up almost ex clusively of porous ZrO2 oxidation protection. From the volatility calculations,  (cr), which does not provide passive  the pre Because oxygen is neither consumed nor produced by Eq.  (5),  the partial pressure of B2O3 (g) does not vary with pO2 (Fig. 1) in this regime. Other gases that form in this regime are listed in  Table III. Pressures for all of  the B-containing gases were de termined in this manner and plotted in Fig. 1. The pressures of  Zr-containing gases were calculated, but not plotted since they  were much lower than those of B species. At the equilibrium pO2 for Eq. (1), the partial pressure of B2O3 (g) above ZrB2 (cr) must be the same as it is above ZrO2 (cr)-B2O3 (l) since the gas is in equilibrium with all of the condensed phases simultaneously.  Pressures for all of the gases met this criterion for consistency,  which can be seen in Fig. 1 for B-containing gases.  The vapor pressures of all of  the gases were calculated to  compile a volatility diagram at 1800 K (Fig. 2(a)). The system  dominant vapor species at 1800 K in air was B2O3 (g) with a pressure of B10\\x002 (344 Pa). Two other species had vapor pressures predicted to be greater than 10\\x0010 (B10\\x005 Pa) at 1800 K in air, BO2 (g) at B10 \\x003 (86 Pa), and BO (g) at B10 \\x008 (B10 \\x003 Pa). Based on the nine order of magnitude increase in the pressure  of the dominant species compared to 1000 K, the rate of B2O3 (l) vaporization would be expected to be significantly higher  evaporation from  at 1800 K, which is  consistent with B2O3 the surface of ZrB2 that is oxidized in air at 1650 K or above. Active oxidation of ZrB2 is also expected at higher temperatures, unless the ZrO2 scale becomes protective. Oxidation of ZrB2-SiC above 2000 K in arc jet testing suggests that ZrO2 may become protective at these temperatures by sintering into a coherent scale.21  Table II.  Volatilization Reactions Involving ZrB2 (cr) as the Primary Condensed Phase  Reactions producing volatile B species  2 O2  ZrB2 (cr)13O2 (g)-ZrO2 (cr)12BO2 (g) ZrB2 ðcrÞþ 5 ðgÞ ! ZrO2 ðcrÞ þ B2O3 ðgÞ ZrB2 (cr)12O2 (g)-ZrO2 (cr)1B2O2 (g) ZrB2 (cr)12O2 (g)-ZrO2 (cr)12BO (g) ZrB2 ðcrÞþ 3 ðgÞ ! ZrO2 ðcrÞ þ B2O ðgÞ ZrB2 (cr)1O2 (g)-ZrO2 (cr)1B2 (g) ZrB2 (cr)1O2 (g)-ZrO2 (cr)12B (g)  2 O2  Reactions producing volatile Zr species  ðcrÞþ 5  2 O2 ðgÞ ! ZrO2 ðgÞ þ B2O3 ZrB2 ZrB2 (cr)12O2 (g)-ZrO (g)1B2O3 (l) ZrB2 ðcrÞ þ 3 2 O2 ðgÞ ! ZrðgÞ þ B2O3 ðlÞ  ðlÞ  Table III.  Volatilization Reactions Involving ZrO2 (cr) and B2O3 (l) as the Primary Condensed Phases  Reactions producing volatile B species  2 O2  B2O3 (l)-B2O3 (g) B2O3 ðlÞ þ 1 ðgÞ ! 2BO2 ðgÞ B2O3 ðlÞ ! B2O2 ðgÞ þ 1 2 O2 ðgÞ B2O3 ðlÞ ! 2BO ðgÞ þ 1 B2O3 (l)-B2O (g)1O2 (g) B2O3 ðlÞ ! B2 2 O2 ðgÞ B2O3 ðlÞ ! 2B ðgÞ þ 3  ðgÞ þ 3  ðgÞ  ðgÞ  2 O2  2 O2  Reactions producing volatile Zr species  ZrO2 (cr)-ZrO2 (g) ZrO2 ðcrÞ ! ZrOðgÞ þ 1 2 O2 ðgÞ ZrO2 (cr)-Zr (g)1O2 (g)  \\x0c', 'IV.  Summary  A volatility diagram was calculated for ZrB2. The diagram defines equilibrium pO2 values for the transition from ZrB2 (cr) to ZrO2 (cr) and B2O3 (l) as a function of temperature. In addition, the diagram shows the vapor pressure of the predominant vapor  species as a function of  temperature and pO2. At 1000 K, exposure to air resulted in passive oxidation due to the formation  of a surface layer of B2O3 (l). The partial pressure of the predominant species, BO2 (g), was predicted to be 10\\x0011 (10\\x006 Pa) at 1000 K in air, consistent with the observation of passive ox idation. Increasing the temperature to 1800 K in air increased  the partial pressure of the predominant B10\\x002 (344 Pa). The substantial consistent with the observed transition from passive to active  species, B2O3 increase in vapor pressure was  (g),  to  oxidation kinetics between these temperatures.  Acknowledgments  The author thanks Drs. Mark Opeka and Inna Talmy of  the Naval Surface  Warfare Center-Carderock Division and Professor Jeff Smith of UMR for many  educational discussions. The SEM image was provided by UMR graduate student  Adam Chamberlain.  References  1R. Telle, L. W. Sigl, and K. Takagi, ‘‘Boride-Based Hard Materials’’; pp. 802-  945 in Handbook of Ceramic Hard Materials, Edited by R. Riedel. Wiley-VCH,  Weinheim, 2000. 2R. A. Cutler, ‘‘Engineering Properties of Borides’’; pp. 787-803 in Ceramics and  Glass: Engineered Materials Handbook, Vol. 4. Edited by S. J. Schneider Jr.. ASM  International, Materials Park, OH, 1991. 3T. A. Jackson, D. R. Eklund, and A. J. Fink,  ‘‘High Speed Propulsion: Per formance Advantage of Advanced Materials,’’ J. Mater. Sci., 39 [19] 5905-13  (2004). 4D. M. Van Wie, D. G. Drewry Jr., D. E. King, and C. M. Hudson,  ‘‘The  Hypersonic Environment: Required Operating Conditions and Design Challeng es,’’ J. Mater. Sci., 39 [19] 5915-24 (2004). 5K. Kuwabara, S. Sakamoto, O. Kida, T. Ishino, T. Kodama, H. Nakajima, T.  Ito, and Y. Hirakawa ‘‘Corrosion Resistance and Electrical Resistivity of ZrB2 Monolithic Refractories’’; pp. 302-5 in Proceedings of UNITECR 2003, the 8th  Biennial Worldwide Conference on Refractories, Osaka,  Japan, October 19-22,  2003, Edited by K. Asano. The Technical Association of Refractories, Tokyo,  2003. 6S. Kinoshita, Y. Miyagishi, and Y. Ono ‘‘Application of Zirconium Diboride  Materials  to Waste Melting Furnace’’; pp. 205-8 in Proceedings of UNITECR  2003,  the 8th Biennial Worldwide Conference on Refractories, Osaka, Japan, Oc tober 19-22, 2003, Edited by K. Asano. The Technical Association of Refractories,  Tokyo, 2003. 7W. C. Tripp and H. C. Graham, ‘‘Thermogravimetric Study of the Oxidation of ZrB2 in the Temperature Range of 8001 to 15001C,’’ J. Electrochem. Soc., 118 [7] 1195-9 (1971). 8L. Kaufman, E. V. Clougherty, and J. B. Berkowitz-Mattuck,  ‘‘Oxidation  Characteristics of Hafnium and Zirconium Diboride,’’ Trans. Metall. Soc. AIME,  239 [4] 458-66 (1967). 9J. B. Berkowitz-Mattuck,  ‘‘High Temperature Oxidation: III. Zirconium and  Hafnium Diborides,’’ J. Electrochem. Soc., 113 [9] 908-14 (1966). 10R. J. Irving and I. G. Worsley,  ‘‘The Oxidation of Titanium Diboride and  Zirconium Diboride at High Temperatures,’’ J. Less-Common Metals, 16, 103-12  (1968). 11J. R. Shappirio, J. J. Finnegan, R. A. Lux, and D. C. Fox,  ‘‘Resistivity, Ox idation Kinetics, and Diffusion Barrier Properties of Thin Film ZrB2,’’ Thin Solid Films, 119 [1] 23-30 (1984). 12H. C. Graham, H. H. Davis, I. V. Kvernes, and W. C. Tripp, ‘‘Microstructural  Features of Oxide Scales Formed on Zirconium Diboride Materials’’; pp. 35-48 in  ZrB2 (cr)  ZrO2 (cr) + B2O3 (l)  BO2 (g)  ZrB2 (cr)  ZrO2 (cr) + B2O3 (l)  B (g)  B (g)  BO (g)  BO (g)  B2O2 (g)  B2O3 (g)  B2O2 (g)  B2O3 (g)  1000 K  1800 K  2500 K  BO2 (g)  (b)  (a)  Fig. 2.  The full volatility diagram for zirconium diboride (ZrB2) at (a) 1800 K and (b) at 1000, 1800, and 2500 K.  Table IV. Partial Pressures of all Vapor Species for Oxidation in Air (pO2 5 0.2) at 1000, 1800, and 2500 K  Species  Partial pressure  1000 K  1800 K  2500 K  BO2 (g) B2O3 (g) B2O2 (g) BO (g)  1 \\x02 10 3 \\x02 10\\x0012 2 \\x02 10\\x0028 3 \\x02 10\\x0022 1 \\x02 10 3 \\x02 10 3 \\x02 10 8 \\x02 10 1 \\x02 10 7 \\x02 10\\x0072  \\x0011  6 \\x02 10 3 \\x02 10\\x003 1 \\x02 10\\x0010 4 \\x02 10\\x008 5 \\x02 10 1 \\x02 10 5 \\x02 10 2 \\x02 10 2 \\x02 10 2 \\x02 10\\x0033  \\x004  2 \\x02 10 3.3 2 \\x02 10\\x004 3 \\x02 10\\x003 1 \\x02 10 4 \\x02 10 1 \\x02 10 4 \\x02 10 5 \\x02 10 7 \\x02 10\\x0019  \\x001  B2O (g) B2 (g) B (g)  \\x0052 \\x0086 \\x0048 \\x0032 \\x0047  \\x0022 \\x0038 \\x0021 \\x0013 \\x0020  \\x0011 \\x0022 \\x0011  ZrO2 (g) ZrO (g)  \\x007  \\x0011  Zr (g)  20 µm  ZrO2  ZrB2  Fig. 3.  A zirconium diboride (ZrB2) ceramic oxidized at 1773 K in air for 30 min showing a surface layer of porous ZrO2.  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Lou, ‘‘Volatility Diagrams for Silica, Silicon Nitride,  and Silicon Carbide and Their Application to High Temperature Decomposition  and Oxidation,’’ J. Am. Ceram. Soc., 73 [10] 2785-3128 (1990).  17J. D. Smith ‘‘Reaction Chemistry and Thermochemistry of Magnesia-Graph ite Systems Containing Anti-Oxidants’’; Ph.D. Thesis, University of Missouri Rolla, 1993. 18B.  Schneider, A. Guette, R. Naslain, M. Cataldi,  and A. Costecalde,  ‘‘A  Theoretical  and  Experimental  Approach  to  the  Active-to-Passive  Transition in the Oxidation of Silicon Carbide,’’ J. Mater. Sci., 33 [2] 535-47  (1998). 19M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1 Hypersonic Aerosurfaces: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 20M. W. Jr. Chase, NIST-JANAF Thermochemical Tables, 4th edition, Amer ican Institute of Physics, Woodbury, NY, 1998. 21A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby.  ‘‘Oxidation of ZrB2-SiC Ceramics Under Atmospheric and Reentry Conditions,’’ Refractory Appl. Trans., 1 [2] 2-8 (2005).  &  \\x0c']"
},{
  "_id": 266,
  "PDF": "Thermal Analysis of Tantalum Carbide-Hafnium Carbide Solid Solutions from Room Temperature to 1400 C.pdf",
  "Text": "['Article  Thermal Analysis of Tantalum Carbide-Hafnium Carbide Solid Solutions from Room Temperature to 1400  C  Cheng Zhang, Archana Loganathan, Benjamin Boesl and Arvind Agarwal *  Plasma Forming Laboratory, Department of Mechanical and Materials Engineering, Florida International  University, Miami, 33139 FL, USA; czhan009@fiu.edu (C.Z.); aloga006@fiu.edu (A.L.); bboesl@fiu.edu (B.B.) * Correspondence: agarwala@fiu.edu; Tel.: +1-305-348-1701  Received: 5 June 2017; Accepted: 25 July 2017; Published: 28 July 2017  Abstract: The thermogravimetric analysis on TaC, HfC, and their solid solutions has been carried out in air up to 1400  C. Three solid solution compositions have been chosen: 80TaC-20 vol % HfC (T80H20), 50TaC-50 vol % HfC (T50H50), and 20TaC-80 vol % HfC (T20H80), in addition to pure  TaC and HfC. Solid solutions exhibit better oxidation resistance than the pure carbides. The onset of oxidation is delayed in solid solutions from 750  C for pure TaC, to 940  C for the T50H50 sample. Moreover, T50H50 samples display the highest resistance to oxidation with the retention of the initial  carbides. The oxide scale formed on the T50H50 sample displays mechanical integrity to prevent  the oxidation of the underlying carbide solid solution. The improved oxidation resistance of the solid solution is attributed to the reaction between Ta2O5 and HfC, which stabilizes the volume changes induced by the formation of Ta2O5 and diminishes the generation of gaseous products. Also, the formation of solid solutions disturbs the atomic arrangement inside the lattice, which delays the  reaction between Ta and O. Both of these mechanisms lead to the improved oxidation resistances of  TaC-HfC solid solutions.  Keywords: tantalum carbide; hafnium carbide; solid solutions; oxidation; thermogravimetric analysis  1. Introduction  The interest in tantalum carbide (TaC) and hafnium carbide (HfC) has been growing in recent years  due to their extremely high melting points, high hardness, and elastic moduli, and more importantly,  their ability to form solid solutions [1-4]. The major applications of these two carbides are leading  edges of reentry vehicles and lining materials for rocket thrusters. In both cases, excellent oxidation resistance is required. However, gaseous products like CO and CO2 are inevitably formed during oxidation, which leads to porous oxide scales that delaminate and spall. The major oxide of TaC is Ta2O5 , which has a melting point of ~1900  C, lower than the desired application temperature of 2000  C or more [5-7]. As a result, the resultant oxide would melt and lose its structural integrity, and fail catastrophically. To reduce the gaseous products as well as retain the integrity of oxide scales under extremely high temperatures, Hafnium diboride (HfB2 ) and its composites have been investigated as promising candidate materials for use on next-generation hypersonic vehicles [8,9]. During oxidation, HfB2 forms a solid scaffold-like structure that mainly consists of HfO2 and molten B2O3 infiltrated between the HfO2 . The resultant oxide scale is dense and crack-free, which provides exceptional the B2O3 starts to evaporate around 700  C and therefore its oxidation resistance. Unfortunately, protection of the underlying materials is lost. The SiC addition was used to stabilize the B2O3 by forming a borosilicate glass phase, which increases the onset evaporation temperature to 1400  C. However, 1400  C is still not high enough to withstand higher application temperatures of 2000  C or more.  Coatings 2017, 7, 111; doi:10.3390/coatings7080111  www.mdpi.com/journal/coatings  coatings\\x0c', 'Coatings 2017, 7, 111  2 of 9  The studies on the solid solutions of TaC-HfC began with the discovery of a TaC0.8HfC0.2 phase that possesses the highest melting point (~4000  C) of known substances [10]. Preliminary oxidation studies have been carried out on TaC0.8HfC0.2 and HfC-rich compositions, but no improvement in the oxidation behavior was observed compared to pure carbides [11-14]. Additionally, sintering aids were  inevitable in those studies, which introduced secondary phases that clouded the understanding of oxidation behaviors. Although TaC and HfC can form solid solutions above 887  C in all compositions, as shown in the phase diagram in Figure 1, oxidation studies on TaC-HfC solid solutions have barely  been investigated.  Figure 1. Phase diagram of TaC and HfC [15]. Copyright 2013 Elsevier.  Recently, Cedillos-Barraza et al. as well as our research group sintered TaC-HfC solid solutions  without sintering additions by spark plasma sintering (SPS) [16,17]. The compositions cover the full spectrum of TaC-HfC solid solutions, and both studies noticed that TaC0.5HfC0.5 has the highest hardness and elastic modulus among the solid solutions. Our group conducted oxidation testing using a plasma jet by exposing these solid solutions and pure carbides to a temperature of ~3000  C at a gas flow rate of sonic speed [18]. In general, the solid solutions showed better oxidation resistance than the pure carbides. The best oxidation resistance was found in the TaC0.5HfC0.5 composition. After 300 s of exposure to such extreme oxidation conditions, the thickness of the oxide scale in TaC0.5HfC0.5 was only 28 µm, which is 1/6 and 1/10 of the oxide scale thickness in pure HfC and TaC, respectively [18]. The improved oxidation mechanism was explained by a newly formed Hf6Ta2O17 phase. More importantly, we found a similar dense solid scaffold and liquid phase structure as reported in HfB2 -SiC/HfB2 -B4C systems that protect the underlying materials [18]. In the case of TaC-HfC solid solutions, the solid scaffold consists of HfO2 and Hf6Ta2O17 , and the liquid phase is made of molten Ta2O5 . Compared to the B2O3 and borosilicate phase in the diboride system, molten Ta2O5 is a much more stable phase with a higher melting point of 1900  C. Hence, the carbide solid solutions exhibit exceptional oxidation resistance.  One question arises after the investigation on the plasma jet oxidation behavior of the carbide solid solutions: How would the carbide solid solutions behave below 1800  C, where the temperature is not high enough to melt the resultant Ta2O5 ? To address this question, we sought to understand the carbide solid solutions from room temperature to 1400  C using the oxidation behavior of thermogravimetric analysis (TGA). Five samples, namely pure TaC, 80TaC-20 vol % HfC (T80H20),  50TaC-50 vol % HfC (T50H50), 20TaC-80 vol % HfC (T20H80), and pure HfC, were chosen. A detailed  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa02\\xa0of\\xa09\\xa0\\xa0The\\xa0studies\\xa0on\\xa0the\\xa0solid\\xa0solutions\\xa0of\\xa0TaC‐HfC\\xa0began\\xa0with\\xa0the\\xa0discovery\\xa0of\\xa0a\\xa0TaC0.8HfC0.2\\xa0phase\\xa0that\\xa0possesses\\xa0the\\xa0highest\\xa0melting\\xa0point\\xa0(~4000\\xa0°C)\\xa0of\\xa0known\\xa0substances\\xa0[10].\\xa0Preliminary\\xa0oxidation\\xa0studies\\xa0have\\xa0been\\xa0carried\\xa0out\\xa0on\\xa0TaC0.8HfC0.2\\xa0and\\xa0HfC‐rich\\xa0compositions,\\xa0but\\xa0no\\xa0improvement\\xa0in\\xa0the\\xa0oxidation\\xa0behavior\\xa0was\\xa0observed\\xa0compared\\xa0to\\xa0pure\\xa0carbides\\xa0[11-14].\\xa0Additionally,\\xa0sintering\\xa0aids\\xa0were\\xa0inevitable\\xa0in\\xa0those\\xa0studies,\\xa0which\\xa0introduced\\xa0secondary\\xa0phases\\xa0that\\xa0clouded\\xa0the\\xa0understanding\\xa0of\\xa0oxidation\\xa0behaviors.\\xa0Although\\xa0TaC\\xa0and\\xa0HfC\\xa0can\\xa0form\\xa0solid\\xa0solutions\\xa0above\\xa0887\\xa0°C\\xa0in\\xa0all\\xa0compositions,\\xa0as\\xa0shown\\xa0in\\xa0the\\xa0phase\\xa0diagram\\xa0in\\xa0Figure\\xa01,\\xa0oxidation\\xa0studies\\xa0on\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0have\\xa0barely\\xa0been\\xa0investigated.\\xa0\\xa0Figure\\xa01.\\xa0Phase\\xa0diagram\\xa0of\\xa0TaC\\xa0and\\xa0HfC\\xa0[15].\\xa0Copyright\\xa02013\\xa0Elsevier.\\xa0Recently,\\xa0Cedillos‐Barraza\\xa0et\\xa0al.\\xa0as\\xa0well\\xa0as\\xa0our\\xa0research\\xa0group\\xa0sintered\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0without\\xa0sintering\\xa0additions\\xa0by\\xa0spark\\xa0plasma\\xa0sintering\\xa0(SPS)\\xa0[16,17].\\xa0The\\xa0compositions\\xa0cover\\xa0the\\xa0full\\xa0spectrum\\xa0of\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0and\\xa0both\\xa0studies\\xa0noticed\\xa0that\\xa0TaC0.5HfC0.5\\xa0has\\xa0the\\xa0highest\\xa0hardness\\xa0and\\xa0elastic\\xa0modulus\\xa0among\\xa0the\\xa0solid\\xa0solutions.\\xa0Our\\xa0group\\xa0conducted\\xa0oxidation\\xa0testing\\xa0using\\xa0a\\xa0plasma\\xa0jet\\xa0by\\xa0exposing\\xa0these\\xa0solid\\xa0solutions\\xa0and\\xa0pure\\xa0carbides\\xa0to\\xa0a\\xa0temperature\\xa0of\\xa0~3000\\xa0°C\\xa0at\\xa0a\\xa0gas\\xa0flow\\xa0rate\\xa0of\\xa0sonic\\xa0speed\\xa0[18].\\xa0In\\xa0general,\\xa0the\\xa0solid\\xa0solutions\\xa0showed\\xa0better\\xa0oxidation\\xa0resistance\\xa0than\\xa0the\\xa0pure\\xa0carbides.\\xa0The\\xa0best\\xa0oxidation\\xa0resistance\\xa0was\\xa0found\\xa0in\\xa0the\\xa0TaC0.5HfC0.5\\xa0composition.\\xa0\\xa0After\\xa0300\\xa0s\\xa0of\\xa0exposure\\xa0to\\xa0such\\xa0extreme\\xa0oxidation\\xa0conditions,\\xa0the\\xa0thickness\\xa0of\\xa0the\\xa0oxide\\xa0scale\\xa0in\\xa0TaC0.5HfC0.5\\xa0was\\xa0only\\xa028\\xa0μm,\\xa0which\\xa0is\\xa01/6\\xa0and\\xa01/10\\xa0of\\xa0the\\xa0oxide\\xa0scale\\xa0thickness\\xa0in\\xa0pure\\xa0HfC\\xa0and\\xa0TaC,\\xa0respectively\\xa0[18].\\xa0The\\xa0improved\\xa0oxidation\\xa0mechanism\\xa0was\\xa0explained\\xa0by\\xa0a\\xa0newly\\xa0formed\\xa0Hf6Ta2O17\\xa0phase.\\xa0More\\xa0importantly,\\xa0we\\xa0found\\xa0a\\xa0similar\\xa0dense\\xa0solid\\xa0scaffold\\xa0and\\xa0liquid\\xa0phase\\xa0structure\\xa0as\\xa0reported\\xa0in\\xa0HfB2‐SiC/HfB2‐B4C\\xa0systems\\xa0that\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0[18].\\xa0In\\xa0the\\xa0case\\xa0of\\xa0\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0the\\xa0solid\\xa0scaffold\\xa0consists\\xa0of\\xa0HfO2\\xa0and\\xa0Hf6Ta2O17,\\xa0and\\xa0the\\xa0liquid\\xa0phase\\xa0is\\xa0made\\xa0of\\xa0molten\\xa0Ta2O5.\\xa0Compared\\xa0to\\xa0the\\xa0B2O3\\xa0and\\xa0borosilicate\\xa0phase\\xa0in\\xa0the\\xa0diboride\\xa0system,\\xa0molten\\xa0Ta2O5\\xa0is\\xa0a\\xa0much\\xa0more\\xa0stable\\xa0phase\\xa0with\\xa0a\\xa0higher\\xa0melting\\xa0point\\xa0of\\xa01900\\xa0°C.\\xa0Hence,\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0exhibit\\xa0exceptional\\xa0oxidation\\xa0resistance.\\xa0One\\xa0question\\xa0arises\\xa0after\\xa0the\\xa0investigation\\xa0on\\xa0the\\xa0plasma\\xa0jet\\xa0oxidation\\xa0behavior\\xa0of\\xa0the\\xa0carbide\\xa0solid\\xa0solutions:\\xa0How\\xa0would\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0behave\\xa0below\\xa01800\\xa0°C,\\xa0where\\xa0the\\xa0temperature\\xa0is\\xa0not\\xa0high\\xa0enough\\xa0to\\xa0melt\\xa0the\\xa0resultant\\xa0Ta2O5?\\xa0To\\xa0address\\xa0this\\xa0question,\\xa0we\\xa0sought\\xa0to\\xa0understand\\xa0the\\xa0oxidation\\xa0behavior\\xa0of\\xa0the\\xa0carbide\\xa0solid\\xa0solutions\\xa0from\\xa0room\\xa0temperature\\xa0to\\xa01400\\xa0°C\\xa0using\\xa0thermogravimetric\\xa0analysis\\xa0(TGA).\\xa0Five\\xa0samples,\\xa0namely\\xa0pure\\xa0TaC,\\xa080TaC‐20\\xa0vol\\xa0%\\xa0HfC\\xa0(T80H20),\\xa050TaC‐50\\xa0vol\\xa0%\\xa0HfC\\xa0(T50H50),\\xa020TaC‐80\\xa0vol\\xa0%\\xa0HfC\\xa0(T20H80),\\xa0and\\xa0pure\\xa0HfC,\\xa0were\\xa0chosen.\\xa0A\\xa0detailed\\xa0\\x0c', 'Coatings 2017, 7, 111  3 of 9  analysis of the oxidation behavior is carried out in the present study using TGA followed by scanning  electron microscopy (SEM).  2. Experimental Details  2.1. Materials  Commercial TaC powder (Inframat Advanced Materials LLC, Manchester, CT, USA) and hafnium  carbide powder (Materion LLC, Cleveland, OH, USA) were used as starting powders. Powders for  solid solutions were mixed by a high-energy vibratory ball milling machine (Across International LLC,  Livingston, NJ, USA) according to their stoichiometric ratio. Pure powders were milled for one hour  separately in a tungsten carbide (WC) jar to breakdown the agglomeration. Subsequently, TaC and  HfC powders were mixed for another hour. The ball to powder ratio was 1:3, using a 6-mm diameter  WC ball. The mixed TaC-HfC powders were consolidated by a spark plasma sintering (SPS) machine (Model 10-4, Thermal Technologies, Santa Rose, CA, USA). The powders were sintered at 1850  C with a heating rate of 100  C/min and a maximum uniaxial pressure of 60 MPa. The holding time was 10 min to ensure the densification. The environment in the vacuum was set at a pressure of 4 Pa.  The details of the processing can be found in our previous study [17].  2.2. TGA Testing and Post-Oxidation Characterization  The TGA oxidation testing was conducted on a small portion (~30 mg) of the sintered pellets.  A thermogravimetric analysis (TGA) analyzer (SDT-Q600, TA Instruments, New Castle, DE, USA)  was used to evaluate the oxidation performance of TaC, HfC, and TaC-HfC solid solution samples. Samples were tested in the air at a heating rate of 5  C/min. The maximum temperature was 1400  C for all samples. The morphologies of the post-oxidation samples were examined by a field emission  SEM (JSM-6330F, JEOL Ltd., Tokyo, Japan).  3. Results and Discussions  3.1. Microstructure and Phases in Sintered TaC, HfC, and TaC-HfC Solid Solutions  The detailed characterization results of  the microstructures  and phases  formed in spark  plasma-sintered TaC-HfC solid solutions were published in our previous paper [17]. For the reader ’s  convenience and the sake of completeness, a summary of the key results is listed in Table 1. All five  samples had high densities varying between 97% and 99%. With the addition of HfC, the samples’  densification increases and the highest densification is found in T20H80 sample. The average grain size  also decreased with the HfC additions. The lattice parameters of the formed solid solutions matched  the theoretical values calculated according to Vegard’s Law. [17]  Table 1. Basic characterizations of the TaC, HfC, and TaC-HfC solid solutions.  Name  Pure TaC T80H20 T50H50 T20H80 Pure HfC  Pellet Density (×103 kg/m3 )  Densification (%)  14.14 13.85 13.26 12.68 12.21  96.7 97.8 98.2 98.8 98.5  Average Grain Size (µm) 6.8 ± 1.4 6.2 ± 2.1 3.8 ± 1.2 3.1 ± 1.1 2.3 ± 0.7  3.2. Macro State Morphology of Post-Oxidation TaC-HfC Solid Solutions  The overall qualitative oxidation resistance of TaC, HfC, and their solid solutions can be inferred  by the morphology of the post-oxidation samples. Appearances of the post-oxidation samples from  the TGA testing are shown in Figure 2. The pure TaC sample showed the worst oxidation resistance,  as it turned into a powdery form with no mechanical integrity (Figure 2a). The oxidized pure HfC,  \\x0c', 'Coatings 2017, 7, 111  4 of 9  on the other hand, displayed a much better oxidation resistance. Structural integrity can still be seen  even though the oxidized sample broke into several pieces (Figure 2b). The post-oxidation samples’  appearances of T80H20 and T20H80 is the combination of oxidation morphology exhibited by pure  TaC and HfC. In the oxidized T80H20 sample (Figure 2c), a large amount of powder is noticed with  a few solid broken pieces. The appearance of the post-oxidation T20H80 (Figure 2d) is analogous to  pure HfC with a small amount of powder. T50H50 is the only sample which does not show significant  spallation and delamination. This suggests that outer layer oxide scale has good mechanical integrity  and can protect the underlying carbide solid solution.  Figure 2. Post-oxidation samples: (a) Pure TaC; (b) Pure HfC; (c) T80H20; (d) T20H80; and (e) T50H50.  3.3. Mass Change during Thermogravimetric Analysis of Carbide Solid Solutions  The weight change curves of TaC-HfC solid solutions are presented in Figure 3 as the degree  of oxidation (α) with the onset oxidation temperatures for five samples. The degree of oxidation (α)  is defined as the ratio of the measured weight change over the theoretical weight change at 100%  conversion. The degree of oxidation (α) is calculated by the following equation:  ∝=  ∆m  ∆m∞  =  mins − mi  mt − mi  where ∆m is the measured weight change, ∆m∞ is the theoretical weight change, mins measured weight, mi is the initial weight, and mt is the theoretical weight. The theoretical weight change is based on the below reactions:  is the real-time  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa04\\xa0of\\xa09\\xa0\\xa0even\\xa0though\\xa0the\\xa0oxidized\\xa0sample\\xa0broke\\xa0into\\xa0several\\xa0pieces\\xa0(Figure\\xa02b).\\xa0The\\xa0post‐oxidation\\xa0samples’\\xa0appearances\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0is\\xa0the\\xa0combination\\xa0of\\xa0oxidation\\xa0morphology\\xa0exhibited\\xa0by\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0In\\xa0the\\xa0oxidized\\xa0T80H20\\xa0sample\\xa0(Figure\\xa02c),\\xa0a\\xa0large\\xa0amount\\xa0of\\xa0powder\\xa0is\\xa0noticed\\xa0with\\xa0a\\xa0few\\xa0solid\\xa0broken\\xa0pieces.\\xa0The\\xa0appearance\\xa0of\\xa0the\\xa0post‐oxidation\\xa0T20H80\\xa0(Figure\\xa02d)\\xa0is\\xa0analogous\\xa0to\\xa0pure\\xa0HfC\\xa0with\\xa0a\\xa0small\\xa0amount\\xa0of\\xa0powder.\\xa0T50H50\\xa0is\\xa0the\\xa0only\\xa0sample\\xa0which\\xa0does\\xa0not\\xa0show\\xa0significant\\xa0spallation\\xa0and\\xa0delamination.\\xa0This\\xa0suggests\\xa0that\\xa0outer\\xa0layer\\xa0oxide\\xa0scale\\xa0has\\xa0good\\xa0mechanical\\xa0integrity\\xa0and\\xa0can\\xa0protect\\xa0the\\xa0underlying\\xa0carbide\\xa0solid\\xa0solution.\\xa0\\xa0Figure\\xa02.\\xa0Post‐oxidation\\xa0samples:\\xa0(a)\\xa0Pure\\xa0TaC;\\xa0(b)\\xa0Pure\\xa0HfC;\\xa0(c)\\xa0T80H20;\\xa0(d)\\xa0T20H80;\\xa0and\\xa0(e)\\xa0T50H50.\\xa03.3.\\xa0Mass\\xa0Change\\xa0during\\xa0Thermogravimetric\\xa0Analysis\\xa0of\\xa0Carbide\\xa0Solid\\xa0Solutions\\xa0The\\xa0weight\\xa0change\\xa0curves\\xa0of\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0are\\xa0presented\\xa0in\\xa0Figure\\xa03\\xa0as\\xa0the\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0with\\xa0the\\xa0onset\\xa0oxidation\\xa0temperatures\\xa0for\\xa0five\\xa0samples.\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0is\\xa0defined\\xa0as\\xa0the\\xa0ratio\\xa0of\\xa0the\\xa0measured\\xa0weight\\xa0change\\xa0over\\xa0the\\xa0theoretical\\xa0weight\\xa0change\\xa0at\\xa0100%\\xa0conversion.\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0(α)\\xa0is\\xa0calculated\\xa0by\\xa0the\\xa0following\\xa0equation:\\xa0∝(cid:3404)∆(cid:1865)∆(cid:1865)(cid:2998)(cid:3404)(cid:1865)(cid:3036)(cid:3041)(cid:3046)(cid:3398)(cid:1865)(cid:3036)(cid:1865)(cid:3047)(cid:3398)(cid:1865)(cid:3036)\\xa0\\xa0where\\xa0∆(cid:1865)\\xa0is\\xa0the\\xa0measured\\xa0weight\\xa0change,\\xa0∆(cid:1865)(cid:2998)\\xa0is\\xa0the\\xa0theoretical\\xa0weight\\xa0change,\\xa0(cid:1865)(cid:3036)(cid:3041)(cid:3046)\\xa0is\\xa0the\\xa0\\xa0real‐time\\xa0measured\\xa0weight,\\xa0(cid:1865)(cid:3036)\\xa0is\\xa0the\\xa0initial\\xa0weight,\\xa0and\\xa0(cid:1865)(cid:3047)\\xa0is\\xa0the\\xa0theoretical\\xa0weight.\\xa0The\\xa0theoretical\\xa0weight\\xa0change\\xa0is\\xa0based\\xa0on\\xa0the\\xa0below\\xa0reactions:\\xa02TaC(cid:3397)92(cid:3415)O(cid:2870)→Ta(cid:2870)O(cid:2873)(cid:3397)2CO(cid:2870)\\xa0(1)\\xa0\\x0c', 'Coatings 2017, 7, 111  9  O2 → Ta2O5 + 2CO2  2TaC +  2 HfC + O2 → HfO2 + CO2  5 of 9  (1)  (2)  Figure 3. TG curve of TaC, HfC, and TaC-HfC solid solutions.  The onset oxidation temperature is defined in Figure 3 when the degree of oxidation increases sharply. Pure TaC starts to oxidize around 750  C, followed by pure HfC which starts around 800  C. The solid solution samples showed improved oxidation resistance, as the oxidation processes have  been delayed compared to the pure carbides, as shown in Figure 2. T80H20 began its oxidation process at 850  C, and T20H80 started around 900  C. Further delay was observed in T50H50, where the onset oxidation temperature was around 940  C. The oxidation process of pure TaC matches the literature description [19,20]. The degree of oxidation increases sharply around 750  C and is completed near 950  C. The transformation from TaC to Ta2O5 involves tremendous volume changes. According to the Pilling-Bedworth (PB) ratio theory, the PB ratio for Ta to Ta2O5 is 2.5 [21]. Such large volume mismatch leads to spallation and delamination. Without any anchoring structure, the resultant oxide will lose its mechanical integrity,  and thus it cannot protect the underlying materials. Consequently, the final product mainly consists of fine powder, as shown in Figure 2a. One spike at 780  C is observed in the TaC oxidation curve in Figure 3. The measured temperature dropped during the oxidation process. The morphology of  the oxidation product suggests that the oxide would delaminate and expose unreacted TaC during  oxidation. The unreacted TaC has a lower temperature than the surface of the delaminated oxide,  so the measured temperature dropped. Figure 3 shows that oxidation process of HfC starts at 800  C. Comparing to the oxidation of pure TaC, the weight increases gradually instead of exhibiting a sharp increase, as evident from the slope.  Such behavior is in accordance with the literature description of the oxidation behavior of HfC [12,22].  HfC has the ability to absorb oxygen without transforming into oxide. The dissolved oxygen sits on the carbon vacancies and forms an oxy-carbide layer. The formed HfO2 , unlike Ta2O5 , does not experience much volume change from the HfC. The PB ratio is only 1.7 [21], which explains the mechanically stable morphology of HfO2 in Figure 2b. The morphologies of the oxidized TaC and HfC are studied by SEM and shown in Figure 4.  The grain size of oxidized TaC is larger than oxidized HfC. Distinct grains can be spotted in the  oxidized HfC sample, but the surface of oxidized TaC is relatively smooth with coalesced grains due  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa05\\xa0of\\xa09\\xa0\\xa0HfC(cid:3397)O(cid:2870)→HfO(cid:2870)(cid:3397)CO(cid:2870)\\xa0(2)\\xa0The\\xa0onset\\xa0oxidation\\xa0temperature\\xa0is\\xa0defined\\xa0in\\xa0Figure\\xa03\\xa0when\\xa0the\\xa0degree\\xa0of\\xa0oxidation\\xa0increases\\xa0sharply.\\xa0Pure\\xa0TaC\\xa0starts\\xa0to\\xa0oxidize\\xa0around\\xa0750\\xa0°C,\\xa0followed\\xa0by\\xa0pure\\xa0HfC\\xa0which\\xa0starts\\xa0around\\xa0800\\xa0°C.\\xa0The\\xa0solid\\xa0solution\\xa0samples\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance,\\xa0as\\xa0the\\xa0oxidation\\xa0processes\\xa0have\\xa0been\\xa0delayed\\xa0compared\\xa0to\\xa0the\\xa0pure\\xa0carbides,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa02.\\xa0T80H20\\xa0began\\xa0its\\xa0oxidation\\xa0process\\xa0at\\xa0850\\xa0°C,\\xa0and\\xa0T20H80\\xa0started\\xa0around\\xa0900\\xa0°C.\\xa0Further\\xa0delay\\xa0was\\xa0observed\\xa0in\\xa0T50H50,\\xa0where\\xa0the\\xa0onset\\xa0oxidation\\xa0temperature\\xa0was\\xa0around\\xa0940\\xa0°C.\\xa0\\xa0Figure\\xa03.\\xa0TG\\xa0curve\\xa0of\\xa0TaC,\\xa0HfC,\\xa0and\\xa0TaC‐HfC\\xa0solid\\xa0solutions.\\xa0The\\xa0oxidation\\xa0process\\xa0of\\xa0pure\\xa0TaC\\xa0matches\\xa0the\\xa0literature\\xa0description\\xa0[19,20].\\xa0The\\xa0degree\\xa0of\\xa0oxidation\\xa0increases\\xa0sharply\\xa0around\\xa0750\\xa0°C\\xa0and\\xa0is\\xa0completed\\xa0near\\xa0950\\xa0°C.\\xa0The\\xa0transformation\\xa0from\\xa0TaC\\xa0to\\xa0Ta2O5\\xa0involves\\xa0tremendous\\xa0volume\\xa0changes.\\xa0According\\xa0to\\xa0the\\xa0Pilling‐Bedworth\\xa0(PB)\\xa0ratio\\xa0theory,\\xa0the\\xa0PB\\xa0ratio\\xa0for\\xa0Ta\\xa0to\\xa0Ta2O5\\xa0is\\xa02.5\\xa0[21].\\xa0Such\\xa0large\\xa0volume\\xa0mismatch\\xa0leads\\xa0to\\xa0spallation\\xa0and\\xa0delamination.\\xa0Without\\xa0any\\xa0anchoring\\xa0structure,\\xa0the\\xa0resultant\\xa0oxide\\xa0will\\xa0lose\\xa0its\\xa0mechanical\\xa0integrity,\\xa0and\\xa0thus\\xa0it\\xa0cannot\\xa0protect\\xa0the\\xa0underlying\\xa0materials.\\xa0Consequently,\\xa0the\\xa0final\\xa0product\\xa0mainly\\xa0consists\\xa0of\\xa0fine\\xa0powder,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa02a.\\xa0One\\xa0spike\\xa0at\\xa0780\\xa0°C\\xa0is\\xa0observed\\xa0in\\xa0the\\xa0TaC\\xa0oxidation\\xa0curve\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0measured\\xa0temperature\\xa0dropped\\xa0during\\xa0the\\xa0oxidation\\xa0process.\\xa0The\\xa0morphology\\xa0of\\xa0the\\xa0oxidation\\xa0product\\xa0suggests\\xa0that\\xa0the\\xa0oxide\\xa0would\\xa0delaminate\\xa0and\\xa0expose\\xa0unreacted\\xa0TaC\\xa0during\\xa0oxidation.\\xa0The\\xa0unreacted\\xa0TaC\\xa0has\\xa0a\\xa0lower\\xa0temperature\\xa0than\\xa0the\\xa0surface\\xa0of\\xa0the\\xa0delaminated\\xa0oxide,\\xa0so\\xa0the\\xa0measured\\xa0temperature\\xa0dropped.\\xa0Figure\\xa03\\xa0shows\\xa0that\\xa0oxidation\\xa0process\\xa0of\\xa0HfC\\xa0starts\\xa0at\\xa0800\\xa0°C.\\xa0Comparing\\xa0to\\xa0the\\xa0oxidation\\xa0of\\xa0pure\\xa0TaC,\\xa0the\\xa0weight\\xa0increases\\xa0gradually\\xa0instead\\xa0of\\xa0exhibiting\\xa0a\\xa0sharp\\xa0increase,\\xa0as\\xa0evident\\xa0from\\xa0the\\xa0slope.\\xa0Such\\xa0behavior\\xa0is\\xa0in\\xa0accordance\\xa0with\\xa0the\\xa0literature\\xa0description\\xa0of\\xa0the\\xa0oxidation\\xa0behavior\\xa0of\\xa0HfC\\xa0[12,22].\\xa0HfC\\xa0has\\xa0the\\xa0ability\\xa0to\\xa0absorb\\xa0oxygen\\xa0without\\xa0transforming\\xa0into\\xa0oxide.\\xa0The\\xa0dissolved\\xa0oxygen\\xa0sits\\xa0on\\xa0the\\xa0carbon\\xa0vacancies\\xa0and\\xa0forms\\xa0an\\xa0oxy‐carbide\\xa0layer.\\xa0The\\xa0formed\\xa0HfO2,\\xa0unlike\\xa0Ta2O5,\\xa0does\\xa0not\\xa0experience\\xa0much\\xa0volume\\xa0change\\xa0from\\xa0the\\xa0HfC.\\xa0The\\xa0PB\\xa0ratio\\xa0is\\xa0only\\xa01.7\\xa0[21],\\xa0which\\xa0explains\\xa0the\\xa0mechanically\\xa0stable\\xa0morphology\\xa0of\\xa0HfO2\\xa0in\\xa0Figure\\xa02b.\\xa0The\\xa0morphologies\\xa0of\\xa0the\\xa0oxidized\\xa0TaC\\xa0and\\xa0HfC\\xa0are\\xa0studied\\xa0by\\xa0SEM\\xa0and\\xa0shown\\xa0in\\xa0Figure\\xa04.\\xa0The\\xa0grain\\xa0size\\xa0of\\xa0oxidized\\xa0TaC\\xa0is\\xa0larger\\xa0than\\xa0oxidized\\xa0HfC.\\xa0Distinct\\xa0grains\\xa0can\\xa0be\\xa0spotted\\xa0in\\xa0the\\xa0oxidized\\xa0HfC\\xa0sample,\\xa0but\\xa0the\\xa0surface\\xa0of\\xa0oxidized\\xa0TaC\\xa0is\\xa0relatively\\xa0smooth\\xa0with\\xa0coalesced\\xa0grains\\xa0due\\xa0to\\xa0localized\\xa0sintering.\\xa0Conventionally,\\xa0sintering\\xa0occurs\\xa0at\\xa00.6\\xa0Tm\\xa0(where\\xa0Tm\\xa0is\\xa0the\\xa0melting\\xa0point).\\xa0The\\xa0melting\\xa0point\\xa0(Tm)\\xa0of\\xa0Ta2O5\\xa0is\\xa01872\\xa0°C;\\xa0hence,\\xa0the\\xa0sintering\\xa0of\\xa0Ta2O5\\xa0will\\xa0start\\xa0around\\xa01250\\xa0°C.\\xa0\\xa0No\\xa0sintering\\xa0is\\xa0expected\\xa0in\\xa0HfO2,\\xa0as\\xa0its\\xa0melting\\xa0point\\xa0is\\xa0around\\xa02800\\xa0°C\\xa0[5].\\xa0The\\xa0pores\\xa0in\\xa0the\\xa0oxidized\\xa0HfC\\xa0are\\xa0from\\xa0the\\xa0formation\\xa0of\\xa0a\\xa0gaseous\\xa0product\\xa0during\\xa0oxidation.\\xa0\\x0c', 'Coatings 2017, 7, 111  6 of 9  to localized sintering. Conventionally, sintering occurs at 0.6 Tm (where Tm is the melting point). The melting point (Tm ) of Ta2O5 is 1872  C; hence, the sintering of Ta2O5 will start around 1250  C. No sintering is expected in HfO2 , as its melting point is around 2800  C [5]. The pores in the oxidized HfC are from the formation of a gaseous product during oxidation.  Figure 4. Post-oxidation SEM images of (a) pure TaC and (b) pure HfC.  Although the post-oxidation samples of T80H20 and T20H80 do not have sufficient mechanical  integrity to protect  the underlying materials  (Figure  2c,d),  these  two solid solutions  showed  improved oxidation resistance by delaying the onset of oxidation temperature, as shown in Figure 3. The oxidation onset temperatures of T80H20 and T20H80 are 850 and 900  C, respectively, which is significantly delayed as compared to pure TaC and HfC. However, both T80H20 and T20H80 still  reached near 100% oxidation. The T50H50 solid solution displayed the best oxidation resistance. Not only did the onset temperature for oxidation increase to near 940  C for the T50H50 samples, only 60% of the oxidation was completed when the temperature reached 1400  C. To further understand the mechanism of improved oxidation resistance in the solid solution  samples, especially in the T50H50 sample,  the morphologies of the post-oxidation samples of the  solid solutions are investigated by SEM and presented in Figure 5. As described earlier, the oxide  morphologies of pure TaC and HfC (Figure 4) are highly porous due to gaseous products resulting  in prominent volume change. The key factor in improving the oxidation resistance of the TaC-HfC solid solutions is to suppress the formation of the Ta2O5 phase. In Figure 5a, the main post-oxidation product has large elongated grains as compared to the grains in Figure 5b,c. The inset shows the top  view of the elongated grains where grains look more equiaxial. The grain enlargement is caused by the dramatic volume change associated with the formation of Ta2O5 and the very high Pilling-Bedworth ratio of 2.5. Additionally, TaC has a cubic crystal structure, whereas Ta2O5 has an orthorhombic crystal structure. The transformation from cubic to orthorhombic leads to elongation in one direction, which is also the reason why Ta2O5 has large grain size. Larger grain size is also possible due to grain growth because of the earlier onset of oxidation in the T80H20 sample as compared to the other two solid  solutions. In the oxidation of the TaC-HfC solid solutions, especially in the T80H20 sample, the formed Ta2O5 can react with the unreacted HfC. The reaction (Reaction (3)) replaces the Ta2O5 with HfO2 , so the volume change during the oxidation process is reduced. The consumption of Ta2O5 reduced the volume change and increased the mechanical integrity of (cid:16)∆G (kJ) = −0.14T − 759.76 the post-oxidation samples in solid solutions.  3Ta2O5 + 7HfC → 7HfO2 + 6TaC + CO  (cid:17)  500-2000  C  T :     [17]  (3)  The oxidation behavior of T20H80 is similar to that of pure HfC, as shown in Figure 5b. Another  beneficial effect of reaction 3 is the delay of the formation of gaseous products in the early oxidation  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa06\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa04.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0pure\\xa0TaC\\xa0and\\xa0(b)\\xa0pure\\xa0HfC.\\xa0Although\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0do\\xa0not\\xa0have\\xa0sufficient\\xa0mechanical\\xa0integrity\\xa0to\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0(Figure\\xa02c,d),\\xa0these\\xa0two\\xa0solid\\xa0solutions\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance\\xa0by\\xa0delaying\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0temperature,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0oxidation\\xa0onset\\xa0temperatures\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0are\\xa0850\\xa0and\\xa0900\\xa0°C,\\xa0respectively,\\xa0which\\xa0is\\xa0significantly\\xa0delayed\\xa0as\\xa0compared\\xa0to\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0However,\\xa0both\\xa0T80H20\\xa0and\\xa0T20H80\\xa0still\\xa0reached\\xa0near\\xa0100%\\xa0oxidation.\\xa0The\\xa0T50H50\\xa0solid\\xa0solution\\xa0displayed\\xa0the\\xa0best\\xa0oxidation\\xa0resistance.\\xa0Not\\xa0only\\xa0did\\xa0the\\xa0onset\\xa0temperature\\xa0for\\xa0oxidation\\xa0increase\\xa0to\\xa0near\\xa0940\\xa0°C\\xa0for\\xa0the\\xa0T50H50\\xa0samples,\\xa0only\\xa060%\\xa0of\\xa0the\\xa0oxidation\\xa0was\\xa0completed\\xa0when\\xa0the\\xa0temperature\\xa0reached\\xa01400\\xa0°C.\\xa0To\\xa0further\\xa0understand\\xa0the\\xa0mechanism\\xa0of\\xa0improved\\xa0oxidation\\xa0resistance\\xa0in\\xa0the\\xa0solid\\xa0solution\\xa0samples,\\xa0especially\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0the\\xa0morphologies\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0investigated\\xa0by\\xa0SEM\\xa0and\\xa0presented\\xa0in\\xa0Figure\\xa05.\\xa0As\\xa0described\\xa0earlier,\\xa0the\\xa0oxide\\xa0morphologies\\xa0of\\xa0pure\\xa0TaC\\xa0and\\xa0HfC\\xa0(Figure\\xa04)\\xa0are\\xa0highly\\xa0porous\\xa0due\\xa0to\\xa0gaseous\\xa0products\\xa0resulting\\xa0in\\xa0prominent\\xa0volume\\xa0change.\\xa0The\\xa0key\\xa0factor\\xa0in\\xa0improving\\xa0the\\xa0oxidation\\xa0resistance\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0is\\xa0to\\xa0suppress\\xa0the\\xa0formation\\xa0of\\xa0the\\xa0Ta2O5\\xa0phase.\\xa0In\\xa0Figure\\xa05a,\\xa0the\\xa0main\\xa0post‐oxidation\\xa0product\\xa0has\\xa0large\\xa0elongated\\xa0grains\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0grains\\xa0in\\xa0Figure\\xa05b,c.\\xa0The\\xa0inset\\xa0shows\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains\\xa0where\\xa0grains\\xa0look\\xa0more\\xa0equiaxial.\\xa0The\\xa0grain\\xa0enlargement\\xa0is\\xa0caused\\xa0by\\xa0the\\xa0dramatic\\xa0volume\\xa0change\\xa0associated\\xa0with\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0and\\xa0the\\xa0very\\xa0high\\xa0Pilling‐Bedworth\\xa0ratio\\xa0of\\xa02.5.\\xa0Additionally,\\xa0TaC\\xa0has\\xa0a\\xa0cubic\\xa0crystal\\xa0structure,\\xa0whereas\\xa0Ta2O5\\xa0has\\xa0an\\xa0orthorhombic\\xa0crystal\\xa0structure.\\xa0The\\xa0transformation\\xa0from\\xa0cubic\\xa0to\\xa0orthorhombic\\xa0leads\\xa0to\\xa0elongation\\xa0in\\xa0one\\xa0direction,\\xa0which\\xa0is\\xa0also\\xa0the\\xa0reason\\xa0why\\xa0Ta2O5\\xa0has\\xa0large\\xa0grain\\xa0size.\\xa0Larger\\xa0grain\\xa0size\\xa0is\\xa0also\\xa0possible\\xa0due\\xa0to\\xa0grain\\xa0growth\\xa0because\\xa0of\\xa0the\\xa0earlier\\xa0onset\\xa0of\\xa0oxidation\\xa0in\\xa0the\\xa0T80H20\\xa0sample\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0two\\xa0solid\\xa0solutions.\\xa0In\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0especially\\xa0in\\xa0the\\xa0T80H20\\xa0sample,\\xa0the\\xa0formed\\xa0Ta2O5\\xa0can\\xa0react\\xa0with\\xa0the\\xa0unreacted\\xa0HfC.\\xa0The\\xa0reaction\\xa0(Reaction\\xa0(3))\\xa0replaces\\xa0the\\xa0Ta2O5\\xa0with\\xa0HfO2,\\xa0so\\xa0the\\xa0volume\\xa0change\\xa0during\\xa0the\\xa0oxidation\\xa0process\\xa0is\\xa0reduced.\\xa0\\xa0Figure\\xa05.\\xa0Cont.\\xa0\\x0c', 'Coatings 2017, 7, 111  7 of 9  stage. Reaction 2 shows that one mole of HfC would react with one mole of oxygen to form one mole of gaseous products (CO or CO2 ). However, in reaction 3, Ta2O5 would consume seven moles of HfC to generate one mole of gaseous product, so the gases produced by HfC can be reduced or at least  delayed. The Gibbs free energy for reaction 3 is computed using Factsage [17]. The delay of the gaseous  product diminishes the cracking within the oxides and enhances the mechanical integrity of the oxide  scales in solid solutions. The well-adhered oxide scale can serve as a protection layer against the  further oxidation of solid solutions. The morphology of the post-oxidation sample of T50H50 is a vivid  proof of this concept, as shown in Figure 5c. The oxide scale of the oxidized T50H50 solid solution is  much denser as compared to the other samples. The oxide grains are mostly equiaxed, suggesting  the moderate volume increase. No large cracks are noticeable in the oxidized T50H50 sample. Thus,  the T50H50 solid solution shows the best oxidation resistance among all the solid solutions.  Figure 5. Post-oxidation SEM images of (a) T80H20 (inset showing the top view of the elongated  grains); (b) T20H80; and (c) T50H50. (Insets are high magnification images).  It must be noted that oxidation onset temperature for T50H50 is 940  C. The onset temperature reveals the interaction between the carbides and oxygen. The study of the absorption of oxygen on  TaC and HfC (001) planes suggests that the oxygen tends to sit on the Hf-C bridge [23]. In the case  of TaC, the preferential oxygen site was on the Ta-Ta bridge. After forming a solid solution, the Ta  atoms are partially replaced by Hf atoms, so the availability of Ta-Ta is disturbed, and the oxidation of the Hf element remains unaffected. Thus, the formation of Ta2O5 is retarded. As discussed earlier, the formation of Ta2O5 is detrimental to oxidation performance, so the solid solutions are expected to have superior oxidation resistances. Among the five samples, abrupt weight increase is observed only  in the pure TaC sample. The weight change in the HfC-contained samples increases steadily, which  corresponds to the adsorption of oxygen, a unique feature of HfC. The solid solutions should have  an abrupt weight increase at a lower temperature similar to pure TaC if the oxidation of TaC has not  been retarded. The maximum delay in the onset of oxidation is found in the T50H50 sample, which is  expected as the Ta-Ta bridges have been disturbed the most by forming solid solutions.  Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa06\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa04.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0pure\\xa0TaC\\xa0and\\xa0(b)\\xa0pure\\xa0HfC.\\xa0Although\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0do\\xa0not\\xa0have\\xa0sufficient\\xa0mechanical\\xa0integrity\\xa0to\\xa0protect\\xa0the\\xa0underlying\\xa0materials\\xa0(Figure\\xa02c,d),\\xa0these\\xa0two\\xa0solid\\xa0solutions\\xa0showed\\xa0improved\\xa0oxidation\\xa0resistance\\xa0by\\xa0delaying\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0temperature,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa03.\\xa0The\\xa0oxidation\\xa0onset\\xa0temperatures\\xa0of\\xa0T80H20\\xa0and\\xa0T20H80\\xa0are\\xa0850\\xa0and\\xa0900\\xa0°C,\\xa0respectively,\\xa0which\\xa0is\\xa0significantly\\xa0delayed\\xa0as\\xa0compared\\xa0to\\xa0pure\\xa0TaC\\xa0and\\xa0HfC.\\xa0However,\\xa0both\\xa0T80H20\\xa0and\\xa0T20H80\\xa0still\\xa0reached\\xa0near\\xa0100%\\xa0oxidation.\\xa0The\\xa0T50H50\\xa0solid\\xa0solution\\xa0displayed\\xa0the\\xa0best\\xa0oxidation\\xa0resistance.\\xa0Not\\xa0only\\xa0did\\xa0the\\xa0onset\\xa0temperature\\xa0for\\xa0oxidation\\xa0increase\\xa0to\\xa0near\\xa0940\\xa0°C\\xa0for\\xa0the\\xa0T50H50\\xa0samples,\\xa0only\\xa060%\\xa0of\\xa0the\\xa0oxidation\\xa0was\\xa0completed\\xa0when\\xa0the\\xa0temperature\\xa0reached\\xa01400\\xa0°C.\\xa0To\\xa0further\\xa0understand\\xa0the\\xa0mechanism\\xa0of\\xa0improved\\xa0oxidation\\xa0resistance\\xa0in\\xa0the\\xa0solid\\xa0solution\\xa0samples,\\xa0especially\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0the\\xa0morphologies\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0of\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0investigated\\xa0by\\xa0SEM\\xa0and\\xa0presented\\xa0in\\xa0Figure\\xa05.\\xa0As\\xa0described\\xa0earlier,\\xa0the\\xa0oxide\\xa0morphologies\\xa0of\\xa0pure\\xa0TaC\\xa0and\\xa0HfC\\xa0(Figure\\xa04)\\xa0are\\xa0highly\\xa0porous\\xa0due\\xa0to\\xa0gaseous\\xa0products\\xa0resulting\\xa0in\\xa0prominent\\xa0volume\\xa0change.\\xa0The\\xa0key\\xa0factor\\xa0in\\xa0improving\\xa0the\\xa0oxidation\\xa0resistance\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions\\xa0is\\xa0to\\xa0suppress\\xa0the\\xa0formation\\xa0of\\xa0the\\xa0Ta2O5\\xa0phase.\\xa0In\\xa0Figure\\xa05a,\\xa0the\\xa0main\\xa0post‐oxidation\\xa0product\\xa0has\\xa0large\\xa0elongated\\xa0grains\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0grains\\xa0in\\xa0Figure\\xa05b,c.\\xa0The\\xa0inset\\xa0shows\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains\\xa0where\\xa0grains\\xa0look\\xa0more\\xa0equiaxial.\\xa0The\\xa0grain\\xa0enlargement\\xa0is\\xa0caused\\xa0by\\xa0the\\xa0dramatic\\xa0volume\\xa0change\\xa0associated\\xa0with\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0and\\xa0the\\xa0very\\xa0high\\xa0Pilling‐Bedworth\\xa0ratio\\xa0of\\xa02.5.\\xa0Additionally,\\xa0TaC\\xa0has\\xa0a\\xa0cubic\\xa0crystal\\xa0structure,\\xa0whereas\\xa0Ta2O5\\xa0has\\xa0an\\xa0orthorhombic\\xa0crystal\\xa0structure.\\xa0The\\xa0transformation\\xa0from\\xa0cubic\\xa0to\\xa0orthorhombic\\xa0leads\\xa0to\\xa0elongation\\xa0in\\xa0one\\xa0direction,\\xa0which\\xa0is\\xa0also\\xa0the\\xa0reason\\xa0why\\xa0Ta2O5\\xa0has\\xa0large\\xa0grain\\xa0size.\\xa0Larger\\xa0grain\\xa0size\\xa0is\\xa0also\\xa0possible\\xa0due\\xa0to\\xa0grain\\xa0growth\\xa0because\\xa0of\\xa0the\\xa0earlier\\xa0onset\\xa0of\\xa0oxidation\\xa0in\\xa0the\\xa0T80H20\\xa0sample\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0two\\xa0solid\\xa0solutions.\\xa0In\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0TaC‐HfC\\xa0solid\\xa0solutions,\\xa0especially\\xa0in\\xa0the\\xa0T80H20\\xa0sample,\\xa0the\\xa0formed\\xa0Ta2O5\\xa0can\\xa0react\\xa0with\\xa0the\\xa0unreacted\\xa0HfC.\\xa0The\\xa0reaction\\xa0(Reaction\\xa0(3))\\xa0replaces\\xa0the\\xa0Ta2O5\\xa0with\\xa0HfO2,\\xa0so\\xa0the\\xa0volume\\xa0change\\xa0during\\xa0the\\xa0oxidation\\xa0process\\xa0is\\xa0reduced.\\xa0\\xa0Figure\\xa05.\\xa0Cont.\\xa0Coatings\\xa02017,\\xa07,\\xa0111\\xa0\\xa07\\xa0of\\xa09\\xa0\\xa0\\xa0Figure\\xa05.\\xa0Post‐oxidation\\xa0SEM\\xa0images\\xa0of\\xa0(a)\\xa0T80H20\\xa0(inset\\xa0showing\\xa0the\\xa0top\\xa0view\\xa0of\\xa0the\\xa0elongated\\xa0grains);\\xa0(b)\\xa0T20H80;\\xa0and\\xa0(c)\\xa0T50H50.\\xa0(Insets\\xa0are\\xa0high\\xa0magnification\\xa0images).\\xa0The\\xa0consumption\\xa0of\\xa0Ta2O5\\xa0reduced\\xa0the\\xa0volume\\xa0change\\xa0and\\xa0increased\\xa0the\\xa0mechanical\\xa0integrity\\xa0of\\xa0the\\xa0post‐oxidation\\xa0samples\\xa0in\\xa0solid\\xa0solutions.\\xa03Ta(cid:2870)O(cid:2873)(cid:3397)7HfC→7HfO(cid:2870)(cid:3397)6TaC(cid:3397)CO\\xa0(ΔG\\xa0(kJ)\\xa0=\\xa0−0.14T\\xa0−\\xa0759.76\\xa0\\xa0\\xa0\\xa0\\xa0\\xa0T:\\xa0500-2000\\xa0°C)\\xa0[17]\\xa0(3)\\xa0The\\xa0oxidation\\xa0behavior\\xa0of\\xa0T20H80\\xa0is\\xa0similar\\xa0to\\xa0that\\xa0of\\xa0pure\\xa0HfC,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa05b.\\xa0Another\\xa0beneficial\\xa0effect\\xa0of\\xa0reaction\\xa03\\xa0is\\xa0the\\xa0delay\\xa0of\\xa0the\\xa0formation\\xa0of\\xa0gaseous\\xa0products\\xa0in\\xa0the\\xa0early\\xa0oxidation\\xa0stage.\\xa0Reaction\\xa02\\xa0shows\\xa0that\\xa0one\\xa0mole\\xa0of\\xa0HfC\\xa0would\\xa0react\\xa0with\\xa0one\\xa0mole\\xa0of\\xa0oxygen\\xa0to\\xa0form\\xa0one\\xa0mole\\xa0of\\xa0gaseous\\xa0products\\xa0(CO\\xa0or\\xa0CO2).\\xa0However,\\xa0in\\xa0reaction\\xa03,\\xa0Ta2O5\\xa0would\\xa0consume\\xa0seven\\xa0moles\\xa0of\\xa0HfC\\xa0to\\xa0generate\\xa0one\\xa0mole\\xa0of\\xa0gaseous\\xa0product,\\xa0so\\xa0the\\xa0gases\\xa0produced\\xa0by\\xa0HfC\\xa0can\\xa0be\\xa0reduced\\xa0or\\xa0at\\xa0least\\xa0delayed.\\xa0The\\xa0Gibbs\\xa0free\\xa0energy\\xa0for\\xa0reaction\\xa03\\xa0is\\xa0computed\\xa0using\\xa0Factsage\\xa0[17].\\xa0The\\xa0delay\\xa0of\\xa0the\\xa0gaseous\\xa0product\\xa0diminishes\\xa0the\\xa0cracking\\xa0within\\xa0the\\xa0oxides\\xa0and\\xa0enhances\\xa0the\\xa0mechanical\\xa0integrity\\xa0of\\xa0the\\xa0oxide\\xa0scales\\xa0in\\xa0solid\\xa0solutions.\\xa0The\\xa0well‐adhered\\xa0oxide\\xa0scale\\xa0can\\xa0serve\\xa0as\\xa0a\\xa0protection\\xa0layer\\xa0against\\xa0the\\xa0further\\xa0oxidation\\xa0of\\xa0solid\\xa0solutions.\\xa0The\\xa0morphology\\xa0of\\xa0the\\xa0post‐oxidation\\xa0sample\\xa0of\\xa0T50H50\\xa0is\\xa0a\\xa0vivid\\xa0proof\\xa0of\\xa0this\\xa0concept,\\xa0as\\xa0shown\\xa0in\\xa0Figure\\xa05c.\\xa0The\\xa0oxide\\xa0scale\\xa0of\\xa0the\\xa0oxidized\\xa0T50H50\\xa0solid\\xa0solution\\xa0is\\xa0much\\xa0denser\\xa0as\\xa0compared\\xa0to\\xa0the\\xa0other\\xa0samples.\\xa0The\\xa0oxide\\xa0grains\\xa0are\\xa0mostly\\xa0equiaxed,\\xa0suggesting\\xa0the\\xa0moderate\\xa0volume\\xa0increase.\\xa0No\\xa0large\\xa0cracks\\xa0are\\xa0noticeable\\xa0in\\xa0the\\xa0oxidized\\xa0T50H50\\xa0sample.\\xa0Thus,\\xa0the\\xa0T50H50\\xa0solid\\xa0solution\\xa0shows\\xa0the\\xa0best\\xa0oxidation\\xa0resistance\\xa0among\\xa0all\\xa0the\\xa0solid\\xa0solutions.\\xa0It\\xa0must\\xa0be\\xa0noted\\xa0that\\xa0oxidation\\xa0onset\\xa0temperature\\xa0for\\xa0T50H50\\xa0is\\xa0940\\xa0°C.\\xa0The\\xa0onset\\xa0temperature\\xa0reveals\\xa0the\\xa0interaction\\xa0between\\xa0the\\xa0carbides\\xa0and\\xa0oxygen.\\xa0The\\xa0study\\xa0of\\xa0the\\xa0absorption\\xa0of\\xa0oxygen\\xa0on\\xa0TaC\\xa0and\\xa0HfC\\xa0(001)\\xa0planes\\xa0suggests\\xa0that\\xa0the\\xa0oxygen\\xa0tends\\xa0to\\xa0sit\\xa0on\\xa0the\\xa0Hf-C\\xa0bridge\\xa0[23].\\xa0In\\xa0the\\xa0case\\xa0of\\xa0TaC,\\xa0the\\xa0preferential\\xa0oxygen\\xa0site\\xa0was\\xa0on\\xa0the\\xa0Ta-Ta\\xa0bridge.\\xa0After\\xa0forming\\xa0a\\xa0solid\\xa0solution,\\xa0the\\xa0Ta\\xa0atoms\\xa0are\\xa0partially\\xa0replaced\\xa0by\\xa0Hf\\xa0atoms,\\xa0so\\xa0the\\xa0availability\\xa0of\\xa0Ta-Ta\\xa0is\\xa0disturbed,\\xa0and\\xa0the\\xa0oxidation\\xa0of\\xa0the\\xa0Hf\\xa0element\\xa0remains\\xa0unaffected.\\xa0Thus,\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0is\\xa0retarded.\\xa0As\\xa0discussed\\xa0earlier,\\xa0the\\xa0formation\\xa0of\\xa0Ta2O5\\xa0is\\xa0detrimental\\xa0to\\xa0oxidation\\xa0performance,\\xa0so\\xa0the\\xa0solid\\xa0solutions\\xa0are\\xa0expected\\xa0to\\xa0have\\xa0superior\\xa0oxidation\\xa0resistances.\\xa0Among\\xa0the\\xa0five\\xa0samples,\\xa0abrupt\\xa0weight\\xa0increase\\xa0is\\xa0observed\\xa0only\\xa0in\\xa0the\\xa0pure\\xa0TaC\\xa0sample.\\xa0The\\xa0weight\\xa0change\\xa0in\\xa0the\\xa0HfC‐contained\\xa0samples\\xa0increases\\xa0steadily,\\xa0which\\xa0corresponds\\xa0to\\xa0the\\xa0adsorption\\xa0of\\xa0oxygen,\\xa0a\\xa0unique\\xa0feature\\xa0of\\xa0HfC.\\xa0The\\xa0solid\\xa0solutions\\xa0should\\xa0have\\xa0an\\xa0abrupt\\xa0weight\\xa0increase\\xa0at\\xa0a\\xa0lower\\xa0temperature\\xa0similar\\xa0to\\xa0pure\\xa0TaC\\xa0if\\xa0the\\xa0oxidation\\xa0of\\xa0TaC\\xa0has\\xa0not\\xa0been\\xa0retarded.\\xa0The\\xa0maximum\\xa0delay\\xa0in\\xa0the\\xa0onset\\xa0of\\xa0oxidation\\xa0is\\xa0found\\xa0in\\xa0the\\xa0T50H50\\xa0sample,\\xa0which\\xa0is\\xa0expected\\xa0as\\xa0the\\xa0Ta-Ta\\xa0bridges\\xa0have\\xa0been\\xa0disturbed\\xa0the\\xa0most\\xa0by\\xa0forming\\xa0solid\\xa0solutions.\\xa0\\xa0\\xa0\\x0c', 'Coatings 2017, 7, 111  4. Conclusions  8 of 9  Through TGA analyses, we investigated the oxidation behavior of pure TaC and HfC as well as  their solid solutions. The solid solutions exhibit improved oxidation resistance compared to the pure  carbides. T50H50 is found to have the best oxidation resistance, followed by T20H80 and T80H20. The onset of oxidation in T50H50 increases by 170 and 120  C as compared to pure TaC and pure HfC, respectively. The improved oxidation resistance can be attributed to the formation of the solid solutions that disturbs the atomic arrangement. Such disturbance delays the formation of Ta2O5 and does not affect the formation of HfO2 . The reaction between Ta2O5 and HfC is also responsible for the superior oxidation performances in the solid solution samples. It diminishes the generation of  gaseous products during oxidation, which reduces the porosity of the oxide scales and leads to the  better protection of the underlying materials. The present study showcases SPS-sintered solid solutions  as a new class of oxidation-resistant materials within ultrahigh temperature ceramics (UHTCs).  Acknowledgments: Cheng Zhang thanks the Florida International University Graduate School for the Dissertation Year Fellowship (DYF) award. Advanced Materials Engineering Research Institute (AMERI), FIU is acknowledged for the research facilities used and the support from its staff in this study.  Author Contributions: Benjamin Boesl and Arvind Agarwal conceived and designed experiments; Archana Loganathan performed the experiments; Cheng Zhang, Archana Loganathan, Benjamin Boesl and Arvind Agarwal analyzed data; Cheng Zhang wrote the paper.  Conflicts of Interest: The authors declare no conflict of interest.  References  1.  2.  3.  4.  5.  Fahrenholtz, W.G.; Wuchina, E.J.; Lee, W.E.; Zhou, Y. Ultra-High Temperature Ceramics: Materials for Extreme  Environment Applications; John Wiley & Sons, Inc.: Hoboken, NJ, USA, 2014.  Louis, L.E. Transition Metal Carbides and Nitrides; Academic Press: New York, NY, USA, 1971.  Upadhya, K.; Yang,  J.; Hoffman, W.P. Materials  for Ultrahigh Temperature Structural Applications.  Am. Ceram. Soc. Bull. 1997, 76, 51-56.  Pierson, H.O. Handbook of Refractory Carbides and Nitrides; William Andrew Publishing: Westwood, NJ, USA,  1996. Opeka, M.M.; Talmy, I.G.; Zaykoski, J.A. Oxidation-based Materials Selection for 2000  C + Hypersonic 2004, 39, 5887-5904.  Aerosurfaces: Theoretical Considerations and Historical Experience.  J. Mater. Sci.  [CrossRef]  6.  Simonenko, E.P.; Sevast’yanov, D.V.; Simonenko, N.P.; Sevast’yanov, V.G.; Kuznetsov, N.T. Promising  Ultra-high Temperature Ceramic Materials for Aerospace Applications. Russ.  J. Inorg. Chem.  2013, 58,  7.  8.  9.  1669-1693. [CrossRef]  Gary, S.P.; Krishnamurthy, N.; Awasthi, A.; Venkatraman, M. The O-Ta (Oxygen-Tantalum) System. J. Phase  Equilib. 1996, 17, 63-77.  Fahrenholtz, W.G.; Hilmas, G.E. Oxidation of Ultra-high Temperature Transition Metal Diboride Ceramics.  Int. Mater. Rev. 2012, 57, 61-72. [CrossRef]  Gasch, M.; Ellerby, D.; Irby, E.; Beckman, S.; Gusman, M.; Johnson, S. Processing, Properties and Arc Jet  Oxidation of Hafnium Diboride/Silicon Carbide Ultra High Temperature Ceramics. J. Mater. Sci. 2004, 39,  5925-5937. [CrossRef]  10.  Agte, C.; Alterhum, H. Investigations of the High-Melting Carbide Systems Connected with Problem of the  Carbon Melting. Z. Technol. Phyzik 1930, 11, 182-191.  11.  Coutright, E.L.; Prater, J.T.; Holcomb, G.R.; Stpierre, G.R.; Rapp, R.A. Oxidation of Hafnium Carbide and  Hafnium Carbide with Additions of Tantalum and Praseodymium. Oxid. Met. 1991, 36, 423-437. [CrossRef]  12.  13.  Ghaffari, S.A.; Faghihi-Sani, M.A.; Golestani-Fard, F.; Ebrahimi, S. Pressureless Sintering of Ta0.8Hf0.2C UHTC in the Presence of MoSi2 . Ceram. Int. 2013, 39, 1985-1989. [CrossRef] Patterson, M.C.L. Advanced HfC-TaC Oxidation Resistance Composite Rocket Thruster. Mater. Manuf.  Process. 1996, 11, 367-379. [CrossRef]  \\x0c', 'Coatings 2017, 7, 111  9 of 9  14.  Rudy, E. Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon System. Part II. Ternary Systems, Vol I.  Ta-HfC-C System; Technical Report: AFML-TR-65-2 Part II Vol. 1; Wright-Patterson Air Force Base, Air Force  Systems Command, Air Force Materials Laboratory: Dayton, OH, USA, 1969.  15.  Ghaffari, S.A.; Faghihi-Sani, M.A.; Golestani-Fard, F.; Nojabayy, M. Diffusion and Solid Solution Formation  Between the Binary Carbides of TaC, HfC, ZrC. Int. J. Refract. Met. Hard Mater. 2013, 41, 180-184. [CrossRef]  16.  Cedillos-Barraza, O.; Grasso, S.; Nasiri, N.A.; Jayaseelan, D.D.; Reece, M.J.; Lee, W.E. Sintering Behavior,  Solid Solution Formation and Characterization of TaC, HfC and TaC-HfC Fabricated by Spark Plasma  Sintering. J. Eur. Ceram. Soc. 2016, 36, 1539-1548. [CrossRef]  17.  Zhang, C.; Gupta, A.; Seal, S.; Boesl, B.; Agarwal, A. Solid Solution Synthesis of Tantalum Carbide-Hafnium  Carbide by Spark Plasma Sintering. J. Am. Ceram. Soc. 2017, 100, 1773-2308. [CrossRef]  18.  Zhang, C. High Temperature Oxidation Study of Tantalum Carbide-Hafnium Carbide Solid Solutions  Synthesized By Spark Plasma Sintering. Ph.D. Thesis, Florida International University, Miami, FL, USA,  18 October 2016.  19.  Desmaison-Brut, M.; Alexandre, N.; Desmaison, J. Comparison of the Oxidation Behavior of Two Dense Hot  Isostatically Pressed Tantalum Carbide (TaC and Ta2C) Materials. [CrossRef]  J. Eur. Ceram. Soc. 1997, 17, 1325-1334.  20.  Zhang, X.; Hilmas, G.E.; Fahrenholtz, W.G. Densification, Mehcanical Properties, and Oxidation Reistance of  TaC-TaB2 Creamics. J. Am. Ceram. Soc. 2008, 91, 4129-4132.  21.  Cramer, S.D.; Covino, B.S., Jr. ASM Handbook Volume 13A: Corrosion: Fundamentals, Testing, and Protection;  ASM International: Geauga County, OH, USA, 2013.  22.  Bargeron, C.B.; Benson, R.C.; Jette, A.N.; Phillips, T.E. Oxidation of Hafnium Carbide in the Temperature Range 1400  to 2060  C. J. Am. Ceram. Soc. 1993, 76, 1040-1046. [CrossRef]  23.  Liu, D.; Deng, J.; Jin, Y.; He, C. Adsorption of Atomic Oxygen on HfC and TaC (110) Surface From Firtst  Principles. Appl. Surf. Sci. 2012, 261, 214-218. [CrossRef]  © 2017 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access  article distributed under the terms and conditions of the Creative Commons Attribution  (CC BY) license (http://creativecommons.org/licenses/by/4.0/).  \\x0c']"
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  "_id": 267,
  "PDF": "Thermal stability of refractory carbide boride composites.pdf",
  "Text": "['Materials Chemistry and Physics 74 (2002) 272-281  Thermal stability of refractory carbide/boride composites  C.R. Wang a,∗  , J.-M. Yang a , W. Hoffman b  a Department of Materials Science and Engineering, University of California, Los Angeles, CA 90095-1595, USA b Air Force Research Laboratory, Edward Air Force Base, Los Angeles, CA 93524-7680, USA  Received 10 May 2001 ; accepted 28 June 2001  Abstract  The thermal stability of refractory carbide/boride composites in oxidizing environment was analyzed. The multi-component thermodynamic stability diagrams for the complex (Hf, Ta, Zr, Si)-C(B)-O composite systems were generated using the proposed linear inequality method. In this method, the stability area of compound is the solution of a set of linear inequalities which is directly obtained from the free energy changes of general chemical reactions. The generated thermodynamic stability diagrams were used to analyze the formation of multilayer oxide scale in HfC-TaC and ZrB 2 -SiC composites, and correlated with experimental results. © 2002 Elsevier Science B.V. All rights reserved.  Keywords: A-B -O ternary systems; HfC-TaC; ZrB  2 -SiC  1.  Introduction                 Advanced materials with temperature capability of over 2000 C are needed for ultra-high temperature structural applications, such as rocket engines and thermal protection systems for space vehicles. The potential ceramic systems that can operate as oxidation-resistant materials with temperature capabilities from 2000 to 2400 C are refractory oxides, carbides, borides and nitrides. There are relatively few refractory oxides that are stable in an oxidizing atmosphere above 2000 C. Hafnia (melting point 2900 C) and zirconia (melting point 2770 C) have sufficiently high melting temperatures and relatively low vapor pressures. However, they undergo solid phase transformations: from monoclinic to tetragonal structure at 1150 and 1650 C and from tetragonal to cubic at 2370 and about 2700 C for zirconia and hafnia, respectively, with a corresponding large volume change. The large volume change would result in the destruction of any large scale component made from these materials, and so in practice the materials must be stabilized with an appropriate additive such as calcium oxide, magnesium oxide, or yttrium oxide. Refractory carbides/borides of Hf, Zr and Ta, etc. are potential candidates for ultra high temperature structural applications since they have melting temperatures considerably higher than their associated oxides, do not undergo any solid phase transformation, and have relatively good thermal shock resistance. However, very limited studies [1 -5]       ∗  Corresponding author.     have been conducted to investigate the oxidation behavior of refractory carbides/borides at temperatures above 2000 C. The oxidation processes of refractory carbides/borides have been shown to be the combined processes of oxygen inward or metal ion outward diffusion and gaseous (or liquid at relatively lower temperatures) by-product outward diffusion through oxide scale [6-15]. Therefore, the oxidation resistances of carbides/borides are mainly influenced by the formation and escaping of gaseous by-products (such as CO, CO2 , B2O3 ) during the oxidation processes which are significantly different from those of their metal counterparts. Typically, the oxide scale formed at high temperature consists of at least two distinctive layers: (1) a much less porous inner oxide layer, and (2) a porous outer oxide layer [12,13]. However, Bargeron et al. [1] indicated that in the (HfO 2−x Cy ) oxidized HfC film, an oxycarbide interlayer was found between the outer porous HfO2 layer and a residual carbide layer with dissolved oxygen in the lattice. The oxide interlayer was found to be a better diffusion barrier for oxygen than either the hafnium oxide or carbide layers. Attempts have also been made to improve the oxidation resistance of refractory carbides/borides with appropriate additives. For example, the addition of SiC can improve the oxidation resistance of both HfB2 and ZrB2 [5]. The effect of adding TaC and PrC2 on the oxidation behavior of HfC has also been evaluated [4]. However, much more experimental work and theoretical study need to be conducted to better understand the oxidation behavior of the refractory carbide/boride composites at ultrahigh temperatures.  0254-0584/02/$ - see front matter  © 2002 Elsevier Science B.V. All  rights reserved.  PII: S 0 2 5 4 0 5 8 4 ( 0 1 ) 0 0 4 8 6 2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  273  (4)  lij = xi zj + yi zj − xj zi − yj zi 2(xi + yi ) where i = 1, . . . , N ; j = 1, . . ., i − 1, i + 1, . . . , N .  According to the thermodynamic properties of equilibria stance i(Ax i By i Ozi ) (where i is the certain value 1 ≤ i ≤ N ) and from the general reactions (1), the stability area of sub(where j = 1, obtained by comparison with other substances j(Axj Byj Ozj ) i − 1, i + 1, . . . , N ) must satisfy the following conditions:  . . . ,  \\x01Gij = \\x01G   where \\x01G   ij     mijG  ij  + RT ln Kpij ≥ 0, − lijG  = G   AxjByjOzj  , and Kpij = aAxjByjOzj  AxiByiOzi     O2  + nijG     A  /P lij  amij  O2  AxiByiOzi  −  − nijG     B  (aA /aB )nij ;  Kpij  aAxiByiOzi  and aAxjByjOzj  is the thermodynamic equilibrium constant for reaction (1), aA and aB are the activities of two metallic components, are the activities of two condensed species considered having the stability areas, PO2 is the oxygen partial pressure, respectively. Assuming that all the activities of condensed species are = 1), and let equal  = aAxjByjOzj  then N − 1 linear inequalities can be written as  kij + lijX1 + nijX2 ≥ 0  is the certain value (1 ≤ i ≤ N ); j = 1,  (5) . . . , i − 1,  The objectives of this research are to study the thermal stability of refractory carbide/boride composite in an oxidizing environment. The multi-component thermodynamic stability diagrams will be constructed using the linear inequality methods for predicting the gas-solid-solid interfacial reactions. The applications of these thermodynamic stability diagrams to interpret the oxidation behavior of several refractory carbide/boride composites will be discussed.  2. Thermodynamic algorithms for the generation of multi-component thermodynamic stability diagrams  Since the potential-pH diagrams were first proposed by Pourbaix [16], the potential-pH diagrams, two partial pressure diagrams, etc. for metal -non-metal-non-metal systems have been widely used in the fields of corrosion science and metallurgy, etc. Methods of calculating thermodynamic equilibria can be categorized into two approaches, namely, the stoichiometric and non-stoichiometric ones [17]. Stoichiometric approach makes use of independent reactions and their equilibrium constants, whereas the non-stoichiometric one is based on linear equations in terms of the chemical potentials of system components to obtain the triple points of the stability areas of three condensed species considered. Yokokawa and co-workers [18 -23] first proposed the use of the generalized chemical potential diagrams for metal-metal-non-metal systems in which log( ordinate was successfully introduced to the generated diagrams by the non-stoichiometric method. The generation of such thermodynamic stability diagram can also be done by stoichiometric method as outlined in this section. The multi-component thermodynamic stability diagrams of metal-metal-non-metal systems (A-B-O) will be selected as an example to describe the construction of a set of linear inequalities. If there are N possible condensed species, then the reactions between condensed species, two metallic components and oxygen can be studied in terms of chemical reactions, which proceed according to the general equation:  aM1 /aM2 ) co mijAxiByiOzi − nijA + nijB + lijO2 = AxjByjOzj  (1)  where Ax i By i Ozi and Axj Byj Ozj are the general forms of condensed species in A-B-O ternary systems, A and B metallic elements, O oxygen, xi , yi , and zj are the stoichiometric numbers of these two compounds, respectively. The elemental balances for each of the three elements, i.e. A, B and O, are then used to find for the above reaction, and the following equations can be yielded, respectively:  nij , mij and lij  zi , xj , yj  mij = xj + yj xi + yi nij = yj xi − xj yi xi + yi  to 1 (i.e. aAxiByiOzi \\x01G   ij  kij =  2.303RT X1 = −log PO2 X2 = +log  aA  aB  where i  i + 1, . . . , N .  Let  Cij = kij Aij = lij Bij = nij  and  Cij−1 = ki  Aij−1 = lij Bij−1 = nij  when i > j  when i < j  then the N − 1 linear inequalities (5) can be rewritten as  AijX1 + BijX2 ≥ Cij  (6)  where j = 1, . . . , N − 1. The set of linear inequalities (6) defines the stability area of condensed compound i (Ax i By i Ozi ). When i changes from 1 to N, the N sets of linear inequalities as (6) can be obtained in the same way. When there are many compounds to be considered in the system, it is necessary to use a computer program to generate the thermodynamic stability diagrams. The computer  (2)  (3)  \\x0c', '274  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  algorithm for the generation of the phase stability diagrams had been proposed by Wang et al. [24], in which a series of relevant objective functions with minimum numbers were introduced to the set of linear inequalities to form the problems of linear programming. The FORTRAN program had been constructed to produce the phase stability diagrams by the use of revised simplex method that was used to solve the problems of linear programming. The slopes of boundary lines can be directly obtained from the specific general reactions (for example, in Eq. (1), the slope of boundary line for the stability areas of two possible condensed species The convex polygon stability area of a certain condensed substance is directly obtained by the solution of a set of linear inequalities. However, unlike the chemical potential method proposed by Yokokawa et al. [18,19] where the fundamental stoichiometric number matrix represents linear equations, the fundamental linear inequalities with at least one coordinate of log(aM1 /aM2 ) in the linear inequality method are directly obtained from the free energy changes of the speci fic general chemical (or electrochemical) reaction equation as shown in (1). Such kinds of specific general chemical (or electrochemical) reaction equations can be obtained in any system which includes at least two metallic components and one non-metallic component.  is nij /lij ).  3. Thermodynamic data  Thermodynamic and physical property data were collected from different sources. However, Thermochemical Properties of Inorganic Substances [26] and JANAF Thermodynamic Tables [27] were used as the primary sources for free-energy values. The standard Gibbs free energies of compounds in (Hf, Ta, Zr, Si)-C(B)-O systems are listed in Tables 1 and 2.  4. Results and discussion  for  log PCO2  log PB2O3 )  The generated diagrams of log(aTa /aHf ) versus log PO2 (log PCO , and for Hf-Ta-C(B)-O systems and log(aSi /aZr ) versus log PO2 (log PB2O3 ) Zr-Si-C(B)-O systems are shown in Figs. 1-7, respectively. From these diagrams, it is very clear that because the oxygen partial pressure through the porous oxide scale remains at relatively high values (linear oxygen pro file) [15], the stability areas of oxides in the oxide scale depend only on the activities of metallic components. The metallic component which has the highest diffusion coefficient is usually forming as the outermost oxide layer. The metallic component which has the lowest diffusion coefficient is usually oxidized to form the inner oxide layer. In the following sections, the thermodynamic stability diagrams will be used to analyze the formation of multilayer oxide scale  Table 1  The standard Gibbs  free energies of compounds considered in (Hf, Ta,  Zr, Si)-C(B)-O systems at 2300 K  Compounds  (g)  O2 CO (g)  (g) (g)a  CO2 B2O3 HfC  b  HfB2 HfO2 Hf (l)  Ta  Ta2C TaC  (l)  Ta2O5 TaB2 Ta2Hf6O19 ZrC  c  Zr  ZrO2 ZrB2 Si (l)  SiO2 SiC  (l)  SiO (g)     −G  (kJ mol  −1 )  599.13  649.30  1013.72  1715.67  451.04  628.27  1451.87  181.78  167.73  594.11  371.71  2834.00  494.50  2973  405.21  173.76  1404.19  581.08  127.15  1180.29  223.25  676.99  a Derived from [25,26]. b Derived from 800-1500 K. c Estimated.  in HfC-TaC and ZrB 2 -SiC composites, and correlated with experimental results.  4.1. Thermal stability and oxidation of HfC-TaC composite  The high temperature oxidation of HfC-TaC composite in the temperature range of 1673-2473 K was studied by Courtright et al. [2] and Patterson et al. [4]. It was found that the oxide layers that formed on HfC-TaC obeyed the parabolic growth kinetics. However, there was a break in the kinetics around 1800 C due to the melting of Ta2O5 . At temperature below 2073 K, the outermost oxide layer     Table 2  The standard Gibbs  free energies of compounds considered in (Zr, Si)-  C(B)-O systems at 1600 K  Compounds  (g)  O2 CO (g)  (g) (g)a  CO2 B2O3 SiC  SiO2 ZrO2 ZrC  ZrSiO4 ZrB2  a Derived from [25,26].     −G  (kJ mol  −1 )  372.14  469.20  799.14  1404.43  167.05  1062.74  1273.02  315.96  2342.70  469.22  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  275  Fig. 1. Thermodynamic stability diagram for Hf-Ta -C -O system at 2300 K with log  P CO = −1 atm at carbide/oxide interfaces.  Fig. 2. Thermodynamic stability diagram for Hf-Ta-B -O system at 2300 K with log  = −1 atm at boride/oxide interfaces.  PB2 O3  \\x0c', '276  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 3. Thermodynamic stability diagram for Hf-Ta -B(C) systems at 2300 K with different  PO2 .  Fig. 4. Thermodynamic stability diagram for Hf-Ta-C -O system at 2300 K with log  = −12 atm.  PO2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  277  Fig. 5. Thermodynamic stability diagram for Hf-Ta-B-O system at 2300 K with log  = −12 atm.  PO2  Fig. 6. Thermodynamic stability diagram for Zr-Si-C(B)-O systems at 2300 K with log  = −12 and log PCO = −0.5 atm.  PO2  \\x0c', '278  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 7. Thermodynamic stability diagram for Zr-Si-B(C) system at 2300 K with log  PCO = log PB2 O3  = −1 atm at carbide (boride)/oxide interfaces.  Fig. 8. Thermodynamic stability diagram for Zr-Si-C(B)-O systems at 1600 K with log  = 0 to −10 atm.  PO2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  279  was found to be Ta2O5 while a mixture of Ta2Hf6O19 and HfO2 was found at temperature above 2073 K. From the generated thermodynamic stability diagram for the HfC-TaC composite as shown in Figs. 1 and 4, the multilayer porous oxide scale would be composed of Ta2O5 (T < 2073 K) or Ta2Hf6O19 (T > 2073 K) as the outer layer and HfO2 as the inner layer. The results obtained from the thermodynamic analysis are in good agreement with the experimental results.  4.2. Thermal stability and oxidation of ZrB2 /SiC composite  The effect of an SiC addition on the oxidation of ZrB2 at temperature up to 1773 K was studied by Tripp et al. [28]. The experimental results showed that when 20 v/o SiC is added to ZrB2 , appreciably more SiO2 glass is formed in the temperature range of 1573-1773 K. The formation of this glass continues to provide oxidation resistance at higher temperatures due to good wettability and surface coverage. The SiO2 -rich glassy layer was found to be the outermost oxide layer, whereas ZrO2 -rich oxide layer was the inner oxide layer. SiC could coexist with ZrO2 in the inner oxide layer.  The generated thermodynamic stability diagrams for SiC/ZrB2 (ZrC) at 1600 K are shown in Figs. 8-10, respectively. From the thermodynamic stability diagrams, it is known that a complex oxide, ZrSiO4 , might be formed during oxidation process. However, the complex oxide ZrSiO4 is not stable when T > 1953 K. At higher temperatures, the oxides which can be formed on the surface are ZrO2 and SiO2 (s, l). Thus, the sequence of the multilayer oxide scale at higher temperatures will be different from that at lower temperatures. The oxide scale in the temperature range of 1573 -1773 K would be SiO 2 (outermost)/ZrSiO4 + SiO2 /ZrSiO4 + ZrO2 /ZrO2 + SiC/ZrB2 + SiC. The multilayer oxide scale in the temperature range of (outermost)/SiO2 (l) + 1973 -2273 C would be SiO2 (l) ZrO2 /ZrO2 + SiC/ZrB2 + SiC. The oxidation behavior of ZrB2 /SiC composite in the temperature range of 2073 -2673 C had been studied by Bull et al. [5]. The experimental results con firmed that silicon oxide was formed as the outermost layer, while zirconium oxide was formed as inner oxide layer. Silicon carbide coexisted with zirconium oxide in the internal oxidation region. It is well known that the formation of liquid glass (such as B2O3 (l) in the temperature range 773-1273 K and SiO 2        Fig. 9. Thermodynamic stability diagram for Zr-Si-C-O system at 1600 K with log  PCO = −1 atm at carbide/oxide interface.  \\x0c', '280  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 10. Thermodynamic stability diagram for Zr-Si-C(B)-O system at 1600 K with log  PCO = log PB2 O3  = −1 atm at boride (carbide)/oxide interfaces.  (l) when T > 1960 K) can separate two totally different gas environments (one with high oxygen partial pressure and another one with very low oxygen partial pressure). The formation of liquid SiO2 (T > 1973 K) might serve as a diffusion barrier. However, at higher temperatures, the volatilization of SiO (g) has signi ficant influence on the stability of liquid SiO2 . The effectiveness of SiC additive on the oxidation resistance of ZrB2 is limited due to the formation of SiO (g). However, at lower temperature (T < 2300 K), glassy SiO2 (l) film is stable and can be used as the diffusion barrier. The existence of such diffusion barrier would greatly improve the oxidation resistance of refractory carbide/boride composites at elevated temperatures.  5. Conclusion  The multi-component thermodynamic stability diagrams for the complex (Hf, Ta, Zr, Si)-C(B)-O composite systems were generated using the proposed linear inequality method. The thermodynamic stability diagrams which include at least two metallic components can clearly represent the stability area of complex oxide in certain log(aM1 /aM2 ) range. The generated thermodynamic stability diagrams  can be used to analyze the formation of multilayer oxide scales during high temperature oxidation of HfC-TaC and ZrB2 -SiC composites.  Acknowledgements  This work was Scienti fic Research.  supported  by  the Air Force Office  of  References  [1] C.B. Bargeron, R.C. Benson, A.N.  Jette,  J.E. Phillips, Oxidation     C, J. Am.  of hafnium carbide in the temperature range 1400-2060  Ceram. Soc. 76 (1993) 1040-1046.  [2] E.L. Courtright, J.T. Prater, G.R. Holcomb, G.R.St. Pierre, R.A. Rapp,  Oxidation of hafnium carbide and hafnium carbide with additions of  tantalum and praseodymium, Oxid. Met. 3 (1991) 423-437.  [3] C.B. Bargeron, R.C. Benson, R.W. Newman, A.N. Jette, T.E. Phillips,  Oxidation mechanisms of hafnium carbide and hafnium diboride in  the temperature range 1400-2100  14 (1993) 29-35.     C, Johns Hopkins APL Tech. Dig.  [4] M.C.L. Patterson, S. He, L.L. Fehrenbacher,  J. Hanigofsky, B.D.  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A A22 (1991) 3075-  3076.  [23] H. Yokokawa, N. Sakai, T. Kawada, M. Dokiya, Chemical potential  diagrams  for  rare-transition metal-oxygen systems.  I. Ln-V-O and  Ln-Mn-O systems, J. Am. Ceram. Soc. 73 (1990) 649-658.  [24] C.R. Wang, Z.B. Zhao,  S.K. Xia, W.Q. Zhang, R.Z. Zhu, The  computer algorithm and program for the generation of phase stability  diagram, CALPHAD 14 (1990) 257-264.  [25] C.M. Chen, K. Aral, A computer program for constructing stability  diagrams  in aqueous  solutions  at  elevated temperatures, Corrosion  38 (1982) 183-190.  [26]  I.  Barin,  O.  Knacke,  Thermochemical  Properties  of  Inorganic  Substances, Springer, Berlin, 1973 and 1977 (Suppl.)  [27] JANAF Thermochemical Tables, 2nd ed., NSRFS-NBS 37, National  Bureau  of  Standards, Washington, DC,  1971  (with  1974-1975  Suppl.).  [28] W.C. Trip, H.H. Davis, H.C. Graham, Effect of an SiC addition on  the oxidation of ZrB2 , Am. Ceram. Soc. Bull. 52 (1973) 612-616.  \\x0c']"
},{
  "_id": 268,
  "PDF": "Thermal stability of refractory carbide-boride composites.pdf",
  "Text": "['Materials Chemistry and Physics 74 (2002) 272-281  Thermal stability of refractory carbide/boride composites  C.R. Wang a,∗  , J.-M. Yang a , W. Hoffman b  a Department of Materials Science and Engineering, University of California, Los Angeles, CA 90095-1595, USA b Air Force Research Laboratory, Edward Air Force Base, Los Angeles, CA 93524-7680, USA  Received 10 May 2001 ; accepted 28 June 2001  Abstract  The thermal stability of refractory carbide/boride composites in oxidizing environment was analyzed. The multi-component thermodynamic stability diagrams for the complex (Hf, Ta, Zr, Si)-C(B)-O composite systems were generated using the proposed linear inequality method. In this method, the stability area of compound is the solution of a set of linear inequalities which is directly obtained from the free energy changes of general chemical reactions. The generated thermodynamic stability diagrams were used to analyze the formation of multilayer oxide scale in HfC-TaC and ZrB 2 -SiC composites, and correlated with experimental results. © 2002 Elsevier Science B.V. All rights reserved.  Keywords: A-B -O ternary systems; HfC-TaC; ZrB  2 -SiC  1.  Introduction                 Advanced materials with temperature capability of over 2000 C are needed for ultra-high temperature structural applications, such as rocket engines and thermal protection systems for space vehicles. The potential ceramic systems that can operate as oxidation-resistant materials with temperature capabilities from 2000 to 2400 C are refractory oxides, carbides, borides and nitrides. There are relatively few refractory oxides that are stable in an oxidizing atmosphere above 2000 C. Hafnia (melting point 2900 C) and zirconia (melting point 2770 C) have sufficiently high melting temperatures and relatively low vapor pressures. However, they undergo solid phase transformations: from monoclinic to tetragonal structure at 1150 and 1650 C and from tetragonal to cubic at 2370 and about 2700 C for zirconia and hafnia, respectively, with a corresponding large volume change. The large volume change would result in the destruction of any large scale component made from these materials, and so in practice the materials must be stabilized with an appropriate additive such as calcium oxide, magnesium oxide, or yttrium oxide. Refractory carbides/borides of Hf, Zr and Ta, etc. are potential candidates for ultra high temperature structural applications since they have melting temperatures considerably higher than their associated oxides, do not undergo any solid phase transformation, and have relatively good thermal shock resistance. However, very limited studies [1 -5]       ∗  Corresponding author.     have been conducted to investigate the oxidation behavior of refractory carbides/borides at temperatures above 2000 C. The oxidation processes of refractory carbides/borides have been shown to be the combined processes of oxygen inward or metal ion outward diffusion and gaseous (or liquid at relatively lower temperatures) by-product outward diffusion through oxide scale [6-15]. Therefore, the oxidation resistances of carbides/borides are mainly influenced by the formation and escaping of gaseous by-products (such as CO, CO2 , B2O3 ) during the oxidation processes which are significantly different from those of their metal counterparts. Typically, the oxide scale formed at high temperature consists of at least two distinctive layers: (1) a much less porous inner oxide layer, and (2) a porous outer oxide layer [12,13]. However, Bargeron et al. [1] indicated that in the (HfO 2−x Cy ) oxidized HfC film, an oxycarbide interlayer was found between the outer porous HfO2 layer and a residual carbide layer with dissolved oxygen in the lattice. The oxide interlayer was found to be a better diffusion barrier for oxygen than either the hafnium oxide or carbide layers. Attempts have also been made to improve the oxidation resistance of refractory carbides/borides with appropriate additives. For example, the addition of SiC can improve the oxidation resistance of both HfB2 and ZrB2 [5]. The effect of adding TaC and PrC2 on the oxidation behavior of HfC has also been evaluated [4]. However, much more experimental work and theoretical study need to be conducted to better understand the oxidation behavior of the refractory carbide/boride composites at ultrahigh temperatures.  0254-0584/02/$ - see front matter  © 2002 Elsevier Science B.V. All  rights reserved.  PII: S 0 2 5 4 0 5 8 4 ( 0 1 ) 0 0 4 8 6 2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  273  (4)  lij = xi zj + yi zj − xj zi − yj zi 2(xi + yi ) where i = 1, . . . , N ; j = 1, . . ., i − 1, i + 1, . . . , N .  According to the thermodynamic properties of equilibria stance i(Ax i By i Ozi ) (where i is the certain value 1 ≤ i ≤ N ) and from the general reactions (1), the stability area of sub(where j = 1, obtained by comparison with other substances j(Axj Byj Ozj ) i − 1, i + 1, . . . , N ) must satisfy the following conditions:  . . . ,  \\x01Gij = \\x01G   where \\x01G   ij     mijG  ij  + RT ln Kpij ≥ 0, − lijG  = G   AxjByjOzj  , and Kpij = aAxjByjOzj  AxiByiOzi     O2  + nijG     A  /P lij  amij  O2  AxiByiOzi  −  − nijG     B  (aA /aB )nij ;  Kpij  aAxiByiOzi  and aAxjByjOzj  is the thermodynamic equilibrium constant for reaction (1), aA and aB are the activities of two metallic components, are the activities of two condensed species considered having the stability areas, PO2 is the oxygen partial pressure, respectively. Assuming that all the activities of condensed species are = 1), and let equal  = aAxjByjOzj  then N − 1 linear inequalities can be written as  kij + lijX1 + nijX2 ≥ 0  is the certain value (1 ≤ i ≤ N ); j = 1,  (5) . . . , i − 1,  The objectives of this research are to study the thermal stability of refractory carbide/boride composite in an oxidizing environment. The multi-component thermodynamic stability diagrams will be constructed using the linear inequality methods for predicting the gas-solid-solid interfacial reactions. The applications of these thermodynamic stability diagrams to interpret the oxidation behavior of several refractory carbide/boride composites will be discussed.  2. Thermodynamic algorithms for the generation of multi-component thermodynamic stability diagrams  Since the potential-pH diagrams were first proposed by Pourbaix [16], the potential-pH diagrams, two partial pressure diagrams, etc. for metal -non-metal-non-metal systems have been widely used in the fields of corrosion science and metallurgy, etc. Methods of calculating thermodynamic equilibria can be categorized into two approaches, namely, the stoichiometric and non-stoichiometric ones [17]. Stoichiometric approach makes use of independent reactions and their equilibrium constants, whereas the non-stoichiometric one is based on linear equations in terms of the chemical potentials of system components to obtain the triple points of the stability areas of three condensed species considered. Yokokawa and co-workers [18 -23] first proposed the use of the generalized chemical potential diagrams for metal-metal-non-metal systems in which log( ordinate was successfully introduced to the generated diagrams by the non-stoichiometric method. The generation of such thermodynamic stability diagram can also be done by stoichiometric method as outlined in this section. The multi-component thermodynamic stability diagrams of metal-metal-non-metal systems (A-B-O) will be selected as an example to describe the construction of a set of linear inequalities. If there are N possible condensed species, then the reactions between condensed species, two metallic components and oxygen can be studied in terms of chemical reactions, which proceed according to the general equation:  aM1 /aM2 ) co mijAxiByiOzi − nijA + nijB + lijO2 = AxjByjOzj  (1)  where Ax i By i Ozi and Axj Byj Ozj are the general forms of condensed species in A-B-O ternary systems, A and B metallic elements, O oxygen, xi , yi , and zj are the stoichiometric numbers of these two compounds, respectively. The elemental balances for each of the three elements, i.e. A, B and O, are then used to find for the above reaction, and the following equations can be yielded, respectively:  nij , mij and lij  zi , xj , yj  mij = xj + yj xi + yi nij = yj xi − xj yi xi + yi  to 1 (i.e. aAxiByiOzi \\x01G   ij  kij =  2.303RT X1 = −log PO2 X2 = +log  aA  aB  where i  i + 1, . . . , N .  Let  Cij = kij Aij = lij Bij = nij  and  Cij−1 = ki  Aij−1 = lij Bij−1 = nij  when i > j  when i < j  then the N − 1 linear inequalities (5) can be rewritten as  AijX1 + BijX2 ≥ Cij  (6)  where j = 1, . . . , N − 1. The set of linear inequalities (6) defines the stability area of condensed compound i (Ax i By i Ozi ). When i changes from 1 to N, the N sets of linear inequalities as (6) can be obtained in the same way. When there are many compounds to be considered in the system, it is necessary to use a computer program to generate the thermodynamic stability diagrams. The computer  (2)  (3)  \\x0c', '274  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  algorithm for the generation of the phase stability diagrams had been proposed by Wang et al. [24], in which a series of relevant objective functions with minimum numbers were introduced to the set of linear inequalities to form the problems of linear programming. The FORTRAN program had been constructed to produce the phase stability diagrams by the use of revised simplex method that was used to solve the problems of linear programming. The slopes of boundary lines can be directly obtained from the specific general reactions (for example, in Eq. (1), the slope of boundary line for the stability areas of two possible condensed species The convex polygon stability area of a certain condensed substance is directly obtained by the solution of a set of linear inequalities. However, unlike the chemical potential method proposed by Yokokawa et al. [18,19] where the fundamental stoichiometric number matrix represents linear equations, the fundamental linear inequalities with at least one coordinate of log(aM1 /aM2 ) in the linear inequality method are directly obtained from the free energy changes of the speci fic general chemical (or electrochemical) reaction equation as shown in (1). Such kinds of specific general chemical (or electrochemical) reaction equations can be obtained in any system which includes at least two metallic components and one non-metallic component.  is nij /lij ).  3. Thermodynamic data  Thermodynamic and physical property data were collected from different sources. However, Thermochemical Properties of Inorganic Substances [26] and JANAF Thermodynamic Tables [27] were used as the primary sources for free-energy values. The standard Gibbs free energies of compounds in (Hf, Ta, Zr, Si)-C(B)-O systems are listed in Tables 1 and 2.  4. Results and discussion  for  log PCO2  log PB2O3 )  The generated diagrams of log(aTa /aHf ) versus log PO2 (log PCO , and for Hf-Ta-C(B)-O systems and log(aSi /aZr ) versus log PO2 (log PB2O3 ) Zr-Si-C(B)-O systems are shown in Figs. 1-7, respectively. From these diagrams, it is very clear that because the oxygen partial pressure through the porous oxide scale remains at relatively high values (linear oxygen pro file) [15], the stability areas of oxides in the oxide scale depend only on the activities of metallic components. The metallic component which has the highest diffusion coefficient is usually forming as the outermost oxide layer. The metallic component which has the lowest diffusion coefficient is usually oxidized to form the inner oxide layer. In the following sections, the thermodynamic stability diagrams will be used to analyze the formation of multilayer oxide scale  Table 1  The standard Gibbs  free energies of compounds considered in (Hf, Ta,  Zr, Si)-C(B)-O systems at 2300 K  Compounds  (g)  O2 CO (g)  (g) (g)a  CO2 B2O3 HfC  b  HfB2 HfO2 Hf (l)  Ta  Ta2C TaC  (l)  Ta2O5 TaB2 Ta2Hf6O19 ZrC  c  Zr  ZrO2 ZrB2 Si (l)  SiO2 SiC  (l)  SiO (g)     −G  (kJ mol  −1 )  599.13  649.30  1013.72  1715.67  451.04  628.27  1451.87  181.78  167.73  594.11  371.71  2834.00  494.50  2973  405.21  173.76  1404.19  581.08  127.15  1180.29  223.25  676.99  a Derived from [25,26]. b Derived from 800-1500 K. c Estimated.  in HfC-TaC and ZrB 2 -SiC composites, and correlated with experimental results.  4.1. Thermal stability and oxidation of HfC-TaC composite  The high temperature oxidation of HfC-TaC composite in the temperature range of 1673-2473 K was studied by Courtright et al. [2] and Patterson et al. [4]. It was found that the oxide layers that formed on HfC-TaC obeyed the parabolic growth kinetics. However, there was a break in the kinetics around 1800 C due to the melting of Ta2O5 . At temperature below 2073 K, the outermost oxide layer     Table 2  The standard Gibbs  free energies of compounds considered in (Zr, Si)-  C(B)-O systems at 1600 K  Compounds  (g)  O2 CO (g)  (g) (g)a  CO2 B2O3 SiC  SiO2 ZrO2 ZrC  ZrSiO4 ZrB2  a Derived from [25,26].     −G  (kJ mol  −1 )  372.14  469.20  799.14  1404.43  167.05  1062.74  1273.02  315.96  2342.70  469.22  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  275  Fig. 1. Thermodynamic stability diagram for Hf-Ta -C -O system at 2300 K with log  P CO = −1 atm at carbide/oxide interfaces.  Fig. 2. Thermodynamic stability diagram for Hf-Ta-B -O system at 2300 K with log  = −1 atm at boride/oxide interfaces.  PB2 O3  \\x0c', '276  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 3. Thermodynamic stability diagram for Hf-Ta -B(C) systems at 2300 K with different  PO2 .  Fig. 4. Thermodynamic stability diagram for Hf-Ta-C -O system at 2300 K with log  = −12 atm.  PO2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  277  Fig. 5. Thermodynamic stability diagram for Hf-Ta-B-O system at 2300 K with log  = −12 atm.  PO2  Fig. 6. Thermodynamic stability diagram for Zr-Si-C(B)-O systems at 2300 K with log  = −12 and log PCO = −0.5 atm.  PO2  \\x0c', '278  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 7. Thermodynamic stability diagram for Zr-Si-B(C) system at 2300 K with log  PCO = log PB2 O3  = −1 atm at carbide (boride)/oxide interfaces.  Fig. 8. Thermodynamic stability diagram for Zr-Si-C(B)-O systems at 1600 K with log  = 0 to −10 atm.  PO2  \\x0c', 'C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  279  was found to be Ta2O5 while a mixture of Ta2Hf6O19 and HfO2 was found at temperature above 2073 K. From the generated thermodynamic stability diagram for the HfC-TaC composite as shown in Figs. 1 and 4, the multilayer porous oxide scale would be composed of Ta2O5 (T < 2073 K) or Ta2Hf6O19 (T > 2073 K) as the outer layer and HfO2 as the inner layer. The results obtained from the thermodynamic analysis are in good agreement with the experimental results.  4.2. Thermal stability and oxidation of ZrB2 /SiC composite  The effect of an SiC addition on the oxidation of ZrB2 at temperature up to 1773 K was studied by Tripp et al. [28]. The experimental results showed that when 20 v/o SiC is added to ZrB2 , appreciably more SiO2 glass is formed in the temperature range of 1573-1773 K. The formation of this glass continues to provide oxidation resistance at higher temperatures due to good wettability and surface coverage. The SiO2 -rich glassy layer was found to be the outermost oxide layer, whereas ZrO2 -rich oxide layer was the inner oxide layer. SiC could coexist with ZrO2 in the inner oxide layer.  The generated thermodynamic stability diagrams for SiC/ZrB2 (ZrC) at 1600 K are shown in Figs. 8-10, respectively. From the thermodynamic stability diagrams, it is known that a complex oxide, ZrSiO4 , might be formed during oxidation process. However, the complex oxide ZrSiO4 is not stable when T > 1953 K. At higher temperatures, the oxides which can be formed on the surface are ZrO2 and SiO2 (s, l). Thus, the sequence of the multilayer oxide scale at higher temperatures will be different from that at lower temperatures. The oxide scale in the temperature range of 1573 -1773 K would be SiO 2 (outermost)/ZrSiO4 + SiO2 /ZrSiO4 + ZrO2 /ZrO2 + SiC/ZrB2 + SiC. The multilayer oxide scale in the temperature range of (outermost)/SiO2 (l) + 1973 -2273 C would be SiO2 (l) ZrO2 /ZrO2 + SiC/ZrB2 + SiC. The oxidation behavior of ZrB2 /SiC composite in the temperature range of 2073 -2673 C had been studied by Bull et al. [5]. The experimental results con firmed that silicon oxide was formed as the outermost layer, while zirconium oxide was formed as inner oxide layer. Silicon carbide coexisted with zirconium oxide in the internal oxidation region. It is well known that the formation of liquid glass (such as B2O3 (l) in the temperature range 773-1273 K and SiO 2        Fig. 9. Thermodynamic stability diagram for Zr-Si-C-O system at 1600 K with log  PCO = −1 atm at carbide/oxide interface.  \\x0c', '280  C.R. Wang et al. / Materials Chemistry and Physics 74 (2002) 272-281  Fig. 10. Thermodynamic stability diagram for Zr-Si-C(B)-O system at 1600 K with log  PCO = log PB2 O3  = −1 atm at boride (carbide)/oxide interfaces.  (l) when T > 1960 K) can separate two totally different gas environments (one with high oxygen partial pressure and another one with very low oxygen partial pressure). The formation of liquid SiO2 (T > 1973 K) might serve as a diffusion barrier. However, at higher temperatures, the volatilization of SiO (g) has signi ficant influence on the stability of liquid SiO2 . The effectiveness of SiC additive on the oxidation resistance of ZrB2 is limited due to the formation of SiO (g). However, at lower temperature (T < 2300 K), glassy SiO2 (l) film is stable and can be used as the diffusion barrier. The existence of such diffusion barrier would greatly improve the oxidation resistance of refractory carbide/boride composites at elevated temperatures.  5. Conclusion  The multi-component thermodynamic stability diagrams for the complex (Hf, Ta, Zr, Si)-C(B)-O composite systems were generated using the proposed linear inequality method. The thermodynamic stability diagrams which include at least two metallic components can clearly represent the stability area of complex oxide in certain log(aM1 /aM2 ) range. The generated thermodynamic stability diagrams  can be used to analyze the formation of multilayer oxide scales during high temperature oxidation of HfC-TaC and ZrB2 -SiC composites.  Acknowledgements  This work was Scienti fic Research.  supported  by  the Air Force Office  of  References  [1] C.B. Bargeron, R.C. Benson, A.N.  Jette,  J.E. Phillips, Oxidation     C, J. Am.  of hafnium carbide in the temperature range 1400-2060  Ceram. Soc. 76 (1993) 1040-1046.  [2] E.L. Courtright, J.T. Prater, G.R. Holcomb, G.R.St. Pierre, R.A. Rapp,  Oxidation of hafnium carbide and hafnium carbide with additions of  tantalum and praseodymium, Oxid. Met. 3 (1991) 423-437.  [3] C.B. Bargeron, R.C. Benson, R.W. Newman, A.N. Jette, T.E. Phillips,  Oxidation mechanisms of hafnium carbide and hafnium diboride in  the temperature range 1400-2100  14 (1993) 29-35.     C, Johns Hopkins APL Tech. Dig.  [4] M.C.L. Patterson, S. He, L.L. Fehrenbacher,  J. Hanigofsky, B.D.  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},{
  "_id": 269,
  "PDF": "Thermo-chemical compatibility of hafnium diboride with yttrium aluminum garnet at 1500°C in air.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  ScienceDirec t  Journal of the European Ceramic Society 35 (2015) 2437-2444  Thermo-chemical compatibility of hafnium diboride with yttrium aluminum garnet at 1500 C in air夽     S.L. Winder a , M.B. Ruggles-Wrenn a,∗  , T. Parthasarathy b , T. Key b , C.M. Carney c  a Air Force Institute of Technology, Wright-Patterson Air Force Base, OH 45433-7765, United States b UES Inc., Dayton, OH 45432, United States c Air Force Research Laboratory, Wright-Patterson Air Force Base, OH 45433-7817, United States  Received 22 December 2014; received in revised form 16 March 2015; accepted 17 March 2015  Available online 6 April 2015  Abstract  Due to its high thermal conductivity and oxidation resistance at very high temperatures, hafnium diboride (HfB2 ) is being considered for use as a leading edge material on sharp-bodied reentry vehicles. In structural applications, HfB2 is likely to operate at elevated temperature in proximity to other refractory materials. The thermo-chemical compatibility of HfB2 with (1) single-crystal Y3Al5O12 (SX YAG) and (2) Al2O3 (alumina) was examined  in multiple contact combinations during furnace heat  treatment exposures at 1500 C  in air. The reaction products were characterized using optical microscopy, SEM, EDS, and XRD. The results are presented and possible mechanisms are discussed. Published by Elsevier Ltd.     Keywords: HfB2 ; YAG; Stability  1.   Introduction  Hafnium diboride  (HfB2 )  and  zirconium diboride  (ZrB2 ) have recently received attention as candidate materials for a variety of aerospace applications, such as sharp  leading edges and thermal protection  systems  for  reusable atmospheric  re-entry vehicles and hypersonic flight vehicles [1-4]. These refractory metal borides, frequently referred  to as ultra-high  temperature ceramics (UHTCs), exhibit a number of unique properties, such as extremely high melting temperatures and hardness, chemical stability, high electrical and thermal conductivity and corrosion resistance. A number of  recent  studies have been devoted  to densification and microstructure of  the UHTCs  [5-10]. Numerous efforts also  focused on  improving  the oxidation  resistance of the UHTCs  [11-15]. The densification and oxidation  studies have been extensive and successful, gaining significant  insight  夽 The views expressed are   those of   the authors and do not reflect   the official  policy or position of the United States Air Force, Department of Defense or the  U.S. Government.  ∗  Corresponding author. Tel.: +1 937 255 3636x4641; fax: +1 937 626 4032.  E-mail address: marina.ruggles-wrenn@afit.edu (M.B. Ruggles-Wrenn).  http://dx.doi.org/10.1016/j.jeurceramsoc.2015.03.022  0955-2219/Published by Elsevier Ltd.  into  the mechanisms and behavior of  the UHTCs as a  function of  temperature, gas chemistry, and pressure. Conversely, the subject of  thermochemical stability of UHTCs with other high-temperature materials has received only  limited attention and most of  that concerned  interactions with additives used  to improve densification and oxidation resistance. Hafnium diboride containing UHTCs are very attractive for high  temperature structural use. Hence  their  thermo-chemical compatibility with other high temperature structural materials is of high  interest.  In particular, YAG  is known  to be  the most creep  resistant oxide  ceramic  and  is  considered  as  a potential matrix and/or fiber material for ceramic matrix composites [16-23]. Mah and Parthasarathy  [24,25] observed an  interaction between HfB2 and  single-crystal Y3Al5O12 (SX YAG). The study focused on degradation of SX YAG  in  the presence of HfB2 , but did not consider  the effects of  the  interaction on HfB2 . Our unpublished work on creep of HfB2 also  revealed interaction when HfB2 was  tested at 1500 C  in air using SX YAG fixtures. In this work, we examine the stability of HfB2 in the presence of SX YAG at 1500 C in air. For comparison, alumina, another high  temperature  structural material  of  choice  is  included in  the  study. Alumina  is  used  extensively  in  oxide matrix              \\x0c', '2438   S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444  Fig. 1. Schematic drawing of   treatment experimental setup: (a) HfB2 in contact with SX YAG, (b) HfB2 and SX YAG separated by Al2O3 spacer, (c) HfB2 in close proximity (3 mm) but not in contact with SX YAG.  the heat   composites  for high  temperature applications. Possible mechanisms  involved  in  the  interaction of HfB2 with SX YAG are discussed.  2. Experimental arrangements                 The HfB2 used  in  this work was fabricated using commercially available HfB2 powder  (Cerac, Milwaukee Wisconsin), which had a purity of 99.5% and a mean particle size of 4.6  \\u242em. A 100 g of HfB2 powder was loaded into a 40-mm graphite die coated with BN and lined with graphite foil. The powder was sintered using spark plasma sintering with a heating and cooling rate of 50 C/min and a maximum temperature of 2100 C. The hold time at 2100 C was 30 min. A pressure of 40 MPa was applied during heating  to 1600 C and held  throughout  the  remainder of the sintering cycle. The pressure was released to 4 MPa during cool-down  to 450 C. Near full density was achieved. The sintered HfB2 billets were sectioned into test samples utilizing electric discharge machining. After machining, all sample surto a 45-\\u242em finish using diamond slurry  faces were polished  to remove surface flaws. The high purity SX YAG rods of 10-mm diameter were supplied by VLOC (New Port Richey, FL). The YAG single crystals were produced using the Czochralski growing process. The 10mm diameter rods were cut from the defect-free portions of the YAG crystals. The 99.8% purity Al2O3 (AD-998) was manufactured by Coorstek (Golden, CO). Hafnium diboride test specimens were exposed to 1500 C in laboratory air for 18 h in different contact combinations with SX YAG and Al2O3 . The schematic drawing of  the heat-treatment experimental setup is shown in Fig. 1. The contact combinations     were:  (1) HfB2 specimen  in contact with SX YAG  (Fig. 1a), (2) HfB2 specimen separated from SX YAG by a 1-mm-thick 99.8% pure Al2O3 spacer  (Fig. 1b), and  (3) HfB2 specimen (3 mm) but not  in close proximity  in contact with SX YAG (Fig. 1c). All heat  treatment experiments were performed  in a horizontal MoSi2 resistance-heated tube furnace. The specimens were placed in a 900-mm long alumina tube of 50-mm diameter closed on both ends with alumina plugs. The  surfaces  of  the  heat-treated  specimens were  examined using X-ray diffraction  (XRD: Rigaku Ultima  IV X-Ray Diffractometer, Tokyo, Japan). The Rietveld refinement program MAUD [26] was used to determine phase fractions from the Xray diffraction patterns. The heat-treated specimens were  then sectioned and polished  to a 1-\\u242em finish using diamond slurry for  examination  and  chemical  analysis. The microstructures were characterized using scanning electron microscopy (SEM: Hitachi TM 3000 Tabletop Microscope, Tokyo, Japan) together with  energy  dispersive  spectroscopy  (EDS: Hitachi Bruker Quantax 70, Tokyo, Japan) for elemental analysis.  3. Results  3.1. Hafnium diboride in contact with SX YAG  Upon completion of  the heat  treatment experiment, a white crustaceous deposit was observed on all surfaces of  the HfB2 specimen indicating an interaction between SX YAG and HfB2 . The deposit appeared  to be  less compact on  the HfB2 surfaces exposed to the furnace environment (Fig. 2a) than on the HfB2 surface in direct contact with SX YAG (Fig. 2b). The  surfaces of  the HfB2 specimen were examined using XRD to determine the composition of the surface scales. Due to the uneven  texture of  the surfaces, parallel beam methodology was employed. Crystallographic Information Files (CIFs) from Crystallography Open Database were used  to characterize  the X-ray returns. The XRD patterns of the HfB2 surfaces revealed the presence of  two phases: cubic yttria-stabilized hafnia  (cYSH) and polycrystalline (PX) YAG. The crustaceous deposit on  the HfB2 surface exposed  to  the  furnace environment was composed of approximately 5% c-YSH and 95% PX YAG (see XRD pattern in Fig. 3).  Fig. 2. Photomicrographs of HfB2 after heat treatment in contact with SX YAG for 18 h at 1500 surface that was in contact with SX YAG.     C in air. (a) HfB2 surface exposed to furnace environment, (b) HfB2  \\x0c', 'S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444   2439  The HfB2 specimen was sectioned and examined with SEM and EDS to further characterize the deposits formed on the HfB2 surfaces exposed to the furnace environment (Fig. 4). The SEM micrograph  in Fig. 4b reveals a presence of  three distinct  layers: unoxidized HfB2 , HfO2 scale, and  the crustaceous surface deposit. The EDS analysis  in Fig. 5 shows  the presence of  the  Fig. 3. HfB2 heat treated in contact with SX YAG for 18 h at 1500 of HfB2 surface exposed to the furnace environment. Cubic yttria-stabilized hafnia (c-YSH) and polycrystalline YAG (Y3Al5O12 ) phases are readily identified.  C in air. XRD     In contrast, the deposit on the HfB2 surface that was in direct 45% contact with SX YAG was composed of approximately  55% PX YAG. The different morphology of  c-YSH and  the surface deposits  in Fig. 2a and b may be caused by  the compositional difference. However,  the more pronounced growth of deposit on  the surface exposed  to  the  furnace environment (Fig. 2a) may be also due  to  lack of physical constraint. The XRD analysis of  the SX YAG surface  that was  in direct contact with HfB2 revealed  the presence of  the following phases: PX YAG, Y4Al2O9 (yttrium aluminum monoclinic, YAM), and YBO3 .     Fig. 4. SEM micrographs of HfB2 after heat treatment in contact with SX YAG for 18 h at 1500 C in air. (a) Cross-sectional view. (b) HfB2 surface exposed to furnace environment. Note the three regions: unoxidized HfB2 , HfO2 scale, and a crustaceous deposit.  Fig. 5. HfB2 after heat treatment in contact with SX YAG for 18 h at 1500 air. The EDS line analysis showing presence of Al, Y, Hf and O in the unoxidized  C in     in   in   HfB2 ,  the HfO2 Y-containing deposit resides on the surface of the oxide scale.  the crustaceous   layer. Note   scale and   that   the Aland  \\x0c', '2440   S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444  Fig. 6. Photomicrographs of HfB2 heat  environment, (b) HfB2 surface that was in contact with Al2O3 .  for 18 h at 1500  treated      C   in air with Al2O3  separating   the HfB2  from SX YAG.   (a) HfB2  surface exposed   to   furnace  Yand Al-containing species in the crustaceous surface deposit, but not in the HfO2 scale.  3.2. Hafnium diboride separated from SX YAG by Al2O3  Profoundly different  results were obtained when  the HfB2 specimen was separated from SX YAG by a 1-mm-thick Al2O3 spacer during heat  treatment. The optical micrographs of  the HfB2 surfaces are shown  in Fig. 6. Note  that  there  is no visible growth of  the crustaceous deposit  seen when HfB2 was heat  treated  in contact with SX YAG  (Fig. 2). The HfB2 surface that was in contact with Al2O3 (Fig. 6b) had a nearly flat, sandy appearance. The surfaces of HfB2 that were exposed  to the furnace environment (Fig. 6a) were also nearly flat, but had a speckled appearance. The XRD pattern  in Fig. 7  shows  that  the deposit on  the HfB2 surface exposed to the furnace environment is comprised 62% monoclinic hafnia (m-HfO2 ) and  38% YBO3 with of  trace amounts of Al2O3 . The HfB2 surface  that was  in direct contact with the alumina spacer was found to be nearly 100% mHfO2 with a slight orientation preference or mostly m-HfO2 with trace amounts of YBO3 . The SEM examination also revealed a porous m-HfO2 scale with a columnar grain structure on all surfaces of  the HfB2 specimen (Fig. 8). Notably,  the oxide scale was much  thicker on  the HfB2 surface  that was  in direct contact with Al2O3 (average scale  thickness   0.19 mm)  than on the HfB2 surfaces exposed to the furnace environment (average scale thickness   0.07 mm).  ≈  ≈  3.3. Hafnium diboride in close proximity but not in contact with SX YAG  Upon completion of  the heat  treatment experiment, a white crustaceous  deposit was  not  observed. The  surfaces  of  the HfB2 specimen were  nearly flat,  cream-colored with  some areas of dark  red  (Fig. 9). The SEM micrographs  in Fig. 10 reveal a difference  in morphology of  the cream-colored and red-hued areas of  the HfB2 surface deposit. The EDS analysis showed  that  the cream-colored areas contained mainly Hf and O, while  the  red-hued areas also had an appreciable Al content. The XRD analysis of  the HfB2 surface determined the presence of m-HfO2 , Al2O3 , YBO3 and a  trace amount of  c-YSH. The HfB2 specimen was sectioned and examined with SEM. The micrographs of the cross-section in Fig. 11 show that a porous m-HfO2 scale with a columnar grain structure has formed on all surfaces of the HfB2 specimen. Remarkably, the m-HfO2 scale on the HfB2 surface facing SX YAG is considerably thicker (0.2 mm on average) than that formed on the HfB2 surface facing away from SX YAG. It is possible that the presence of SX YAG accelerates the oxidation of HfB2 .  4. Discussion     To summarize the results, heat treatment of HfB2 at 1500 C in air  in  the presence of SX YAG caused formation of Yand Al-containing deposits on the surface of HfB2 . Heat treatment of HfB2 in contact with YAG resulted in a surface deposit consisting of c-YSH and PX YAG. Conversely, heat treatment of HfB2 separated from SX YAG by Al2O3 gave rise  to a surface  layer of HfB2 in close proximity (3 mm) but not in contact with SX consisting of m-HfO2 , YBO3 and Al2O3 . Finally, heat treatment YAG produced a surface deposit consisting of m-HfO2 , YBO3 , Al2O3 and a trace amount of c-YSH. A summary of the deposits produced on the surface of HfB2 samples is given in Fig. 12.  treated for 18 h at 1500  Fig. 7. HfB2 heat  in air with Al2O3 separating  HfB2 from SX YAG. XRD pattern of  the HfB2 surface exposed  to  the furnace environment shows m-HfO2 , YBO3 and Al2O3 phases.  the     C   \\x0c', 'S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444   2441  Fig. 9. Photomicrograph of HfB2 heat treated for 18 h at 1500 proximity (3 mm) but not in contact with SX YAG.     C in air in close  It  is believed  that  the  formation and  transport of gaseous species  is  required  for  the aforementioned surface deposits  to occur. Mah et al.  reported  the carbothermal  reduction of SX YAG  into YAlO3 (yttrium aluminum perovskite, YAP), YAM, and Y2O3 in the presence of CO (g) due to loss of Al-containing gases [27]. Mah and Parthasarathy proposed that a reduction of SX YAG  into AlBO2 (g) and YAP  in  the presence of B2O was thermodynamically extremely favorable at 1600 C  in air [24]. Recent studies demonstrated  that B2O3 (g)  reacted with Y3+ cations in various compositions to produce YBO2 (g) [28,29]. The  intermediate gaseous species  is B2O3 (g) produced  in the course of oxidation of HfB2 according to [14,15,30-32]:     HfB2 (s)   O2 →  + 5  2   HfO2 (s)   +  B2O3 (g)   (1)  We postulate that in all heat treatment experiments reported here, the B2O3 (g) reacts with SX YAG to produce YAM, AlBO2 (g) and YBO2 (g) according to:  + 5  B2O3 (g) → 1 2  Y4Al2O9 (s) +   4AlBO2 (g)  Y3Al5O12 (s)   2  YBO2 (g) + 5 2  +  O2 (g)  (2)  Fig. 8. SEM micrographs of HfB2 heat  Al2O3 separating  the HfB2 from SX YAG. (a) Cross-sectional view. (b) HfB2 surface exposed  to furnace environment. (c) HfB2 surface  that was  in contact  treated for 18 h at 1500  C   in air with     with Al2O3 .  When HfB2 was heat treated in contact with SX YAG, c-YSH containing crustaceous deposit was found on the surface of the m-HfO2 scale. We postulate  that  the surface m-HfO2 is  transformed  to c-YSH due  to Y3+ cations replacing Hf4+ cations  in  Fig. 10. HfB2 specimen heat treated for 18 h at 1500 (a) red-hued and (b) cream-colored areas.     C in air in close proximity (3 mm) but not in contact with SX YAG. SEM micrographs of the HfB2 surface,    \\x0c', '2442   S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444     the lattice structure. It has been shown that at 1500 C approximately 4-31 mol% Y3+ cations is required to transform m-HfO2 to c-YSH [33,34]. In this case, possible sources of Y3+ cations are:  (i)  the YBO2 gas produced  in  reaction  (2) and adsorbed onto  the surface of  the m-HfO2 scale and (ii)  the diffusion of Y3+ cations from SX YAG  into  the surface of  the scale. Diffusion of Y3+ cations, which proceeds through the least restrictive structural regions (i.e. a free surface)  is further assisted by  the the Hf4+ cation vacancies.  presence of  In a  recent study, Xia noted that Hf4+ cation vacancies within the m-HfO2 scale tend to migrate  towards  the surface  [35]. Hence,  transformation of m-HfO2 to c-YSH  takes place on  the surface of  the m-HfO2 scale. It  is also believed  that as  the AlBO2 and YBO2 produced in reaction (2) are adsorbed onto c-YSH surface, molecules of both gases are dissociated due to a valence electron imbalance. The resulting molecular fragments react  to form PX YAG and a gaseous boron-containing species (assumed to be B2O) on all surfaces of  the HfB2 specimen heat  treated  in contact with SX YAG.  3YBO2 (g) +   5AlBO2 (g) →   Y3Al5O12 (s) +   4B2O (g)  (3)     The  free energies of  reactions  (2) and  (3) were evaluated using available data  [36,37]. The  reactions  (2) and  (3) were thermodynamically highly favorable at 1500 C. When HfB2 is separated from SX YAG by Al2O3 during heat treatment,  the AlBO2 and YBO2 gases produced  in  reaction (2) are adsorbed onto  the  surface of  the m-HfO2 scale,  then transition into Al2O3 and YBO3 . An AlBO2 molecule consists of an Al1+ cation and a (BO2 )1− anion that form a single bond with the bond enthalpy of 502 kJ/mol [38]. The Y1+ cation single bond the (BO2 )1− with  has an enthalpy of 714 kJ/mol [38]. The mHfO2 acts as a catalyst providing sufficient energy to dissociate AlBO2 to form Al2O3 and a boron-containing gas. When YBO2 gas  is formed, a Y atom adopts a 1+ oxidation state  to balance the ionic charge of (BO2 )1− . However, because Y has a standard oxidation state of 3+, yttrium borate is likely to be more stable in a solid form as YBO3 . Therefore, it is likely that when YBO2 gas is adsorbed onto  the surface of m-HfO2 it  transitions  to YBO3 solid  to achieve a  lower energy  level. It  is recognized  that mHfO2 is nonstoichiometric [39]. Furthermore, m-HfO2 has been shown to lose O2− anions [39], which are readily incorporated by YBO2 to achieve a more stable state as YBO3 . When HfB2 specimen was heat  treated  in close proximity but not in contact with SX YAG, Al2O3 and YBO3 were found on the surface of HfB2 . The AlBO2 and YBO2 gases produced in reaction (2) interact with the m-HfO2 scale and transition to Al2O3 and YBO3 as was  the case when HfB2 was separated from SX YAG by Al2O3 during heat treatment. It is proposed that the presence of a continuous c-YSH layer formed when HfB2 is heat  treated  in contact with SX YAG is  required  to  facilitate  formation of PX YAG. Recent studies reported  that cubic  form of yttria-stabilized zirconia  (c-YSZ) exhibits greater catalytic properties and dissociates molecules of various gases more readily  than monoclinic zirconia (m-ZrO2 ) [40-44]. Because hafnia closely resembles zirconia in chemical  Fig. 11. SEM micrographs of HfB2 heat treated for 18 h at 1500 C in air in close proximity (3 mm) but not in contact with SX YAG. (a) Cross-sectional view. (b) HfB2 surface facing SX YAG. (c) HfB2 surface facing away from SX YAG.     properties  [45-48],  it  is  likely  that  the same conclusion holds regarding the catalytic properties of c-YSH vs. those of m-HfO2 . As Y3+ cations replace Hf4+ cations  in  the  lattice structure of m-HfO2 , an O2− anion is removed from the lattice. An increase in oxygen vacancies has been shown  to  increases  the catalytic capability, including adsorbtivity, of a material [41,43,44,49]. It is believed  that  the c-YSH acts  to dissociate  the molecules of both AlBO2 and YBO2 gases and promote reaction (3) to form PX YAG.  \\x0c', 'S.L. Winder et al. / Journal of the European Ceramic Society 35 (2015) 2437-2444   2443  Fig. 12. Schematic showing surface deposits produced during heat   treatment  for 18 h at 1500 C in air: (a) HfB2 in contact with SX YAG, (b) HfB2 and SX YAG separated by Al2O3 spacer, (c) HfB2 in close proximity (3 mm) but not in contact with SX YAG.     5. Summary and conclusions  A  study of  thermo-chemical  compatibility of HfB2 with SX YAG was  carried  out  at  1500 C  in  air. Yttrium and Al-containing deposits were found on the surface of HfB2 specimens heat treated in contact with or in the vicinity of SX YAG, indicating that SX YAG is not stable in the presence of HfB2 at 1500 C in air. It is proposed that SX YAG reacts with B2O3 gas produced during oxidation of HfB2 and  is reduced  into YAM, AlBO2 (g) and YBO2 (g). The gases are adsorbed onto  the HfO2 scale; thus, forming various Yand Al-containing deposits. Surface properties of  the HfO2 scale determine  the nature and composition of the surface deposits. Mechanisms for formation of these deposits were proposed. The  results of  the current study demonstrate  that SX YAG undergoes degradation in the presence of HfB2 at elevated temperatures  in an oxidizing environment. Because both yttrium aluminates and transition metal diborides are being considered for high  temperature applications, environmental stabilities of probable combinations of  these materials must be  investigated further.        6. Summary of novel conclusions  A  study of  thermo-chemical  compatibility of HfB2 with SX YAG was  carried  out  at  1500 C  in  air. Yttrium and Al-containing deposits were found on the surface of HfB2 specimens heat treated in contact with or in the vicinity of SX YAG, indicating that SX YAG is not stable in the presence of HfB2 at 1500 C in air. It is proposed that SX YAG reacts with B2O3 gas produced during oxidation of HfB2 and  is reduced  into YAM, AlBO2 (g) and YBO2 (g). The gases are adsorbed onto  the HfO2 scale; thus, forming various Yand Al-containing deposits. Surface properties of  the HfO2 scale determine  the nature and composition of the surface deposits. Mechanisms for formation of these deposits were proposed.        Acknowledgment  The support of  the Air Force Office of Scientific Research (Grant F1ATA03039J001), Dr. Ali Sayir, Program Director,  is highly appreciated.  References  [1] Levine SR, Opila EJ, Halbig MC, Kiser JD, Singh M, Salem JA. Evaluation  of ultra-high temperature ceramics for aeropropulsion use. J Eur Ceram Soc  2002;22:2757-67.  [2] Wuchina E, Opeka M, Causey S, Buesking K, Spain   J, Cull A, et al.  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},{
  "_id": 270,
  "PDF": "Thermochemical and Mechanical Stabilities of the Oxide Scale of ZrB21SiC and Oxygen Transport Mechanisms.pdf",
  "Text": "['Thermochemical and Mechanical Stabilities of the Oxide Scale of  ZrB21SiC and Oxygen Transport Mechanisms  Ju Li and Thomas J. Lenosky  w  Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia, Pennsylvania 19104  Clemens J. Fo¨ rst and Sidney Yip  Department of Nuclear Science and Engineering and Department of Materials Science and Engineering, Massachusetts  Institute of Technology, Cambridge, Massachusetts 02139  Refractory diboride with silicon carbide additive has a unique  oxide scale microstructure with two condensed oxide phases (solid1liquid), and demonstrates oxidation resistance superior  to either monolithic diboride or silicon carbide. We rationalize  that this is because the silica-rich liquid phase can retreat out ward to remove the high SiO gas volatility region, while still  holding onto the  zirconia skeleton mechanically by capillary  forces,  to form a ‘‘solid pillars,  liquid roof ’’ scale architecture  and maintain barrier function. Basic assessment of the oxygen  carriers  in the borosilicate  liquid in oxygen-rich condition is  performed  using  first-principles  calculations.  It  is  estimated  from entropy and mobility arguments that above a critical temperature TCB15001C, the dominant oxygen carriers should be network defects, such as peroxyl linkage or oxygen-deficient \\x03 as in the Deal-Grove model. instead of molecular O2 These network defects will lead to sublinear dependence of the  centers,  oxidation rate with external oxygen partial pressure. The pres ent work suggests  that  there could be significant room in im proving the high-temperature oxidation resistance by refining  the oxide scale microstructure as well as controlling the glass  chemistry.  I.  Introduction  REFRACTORY diborides (HfB2, ZrB2) with 20-30 vol% SiC additive are prominent ultrahigh-temperature ceramics withstanding temperatures 2000 K and above.1,2 During oper ation in air  its  surface is oxidized, giving rise to a crystalline  oxide skeleton (HfO2, ZrO2) and a silica-rich borosilicate liquid that wets it,3-10 produced by the reactions:  ZrB2 ðcÞ þ ð5=2ÞO2 ! ZrO2 ðcÞ þ B2O3 ðl ; gÞ  (1)  SiCðcÞ þ ð3=2ÞO2 ! SiO2 ðl Þ þ COðgÞ  (2)  respectively.  Intense research is ongoing to characterize scale as a barrier against oxygen11-23  and  enhance  this  (the  scale  microstructure can be seen in, for example, Fig. 4 of Opila et al.15), which apparently is superior to that of either monolithic  diboride or SiC at the intended high temperatures (see Fig. 1). The reason for the first superiority (ZrB21SiC4ZrB2) is well understood. Pure B2O3 melts at 4501C and evaporates rather  quickly above 11001C.  In contrast, SiO2 former (pure SiO2 has a glass transition temperature of 11751C), with much larger viscosity as well as much smaller  is a strong network  equilibrium vapor pressure than B2O3 (see Figs. 15 and 16 of Monteverde and Bellosi).12,13 Thus, oxygen diffusion should be  more  sluggish in the  silica-rich liquid than in pure B2O3 which furthermore will be evaporating rather quickly above 11001C.  (l),  The reason for  the second superiority (ZrB21SiC4SiC) less well understood. The crystalline oxide phase ZrO2(c) formed is often highly porous—although in arc jet testing above 2000 K  is  it appears that ZrO2(c) could sinter into a less porous compact layer,16,17 thereby potentially becoming protective also. Whether  the in situ formed ZrO2(c), when fully dense, barrier to oxygen diffusion as SiO2 (l) in open-circuit condition, is an interesting question that depends on its electronic conduc is as a significant  tivity, which in turn depends on how the  charge defects are  compensated inside the crystal, related to the amount of impu rities. Irrespective of  the outcome of  that discussion, however,  if ZrO2(c) is quite porous its barrier function is lost, because gasphase diffusion through the percolating cracks and pores, even  in the Knudsen diffusion regime, is much easier than diffusion in  condensed phases. In that case then, properties of the silica-rich  liquid will control the effective barrier function of the scale, be cause it flows to fill in the cracks and pores of ZrO2(c), as well as forming an overlayer on top (a ‘‘liquid roof,’’ see for instance Fig. 4 of Opila et al.15 and Fig. 6 of Rezaie et al.22),  thus  occupying both parallel and serial oxygen transport routes.  We suggest an (ZrB21SiC4SiC) protective condensed-phase oxide  explanation  for  the  second  superiority  in Section II, based on the notion that a  scale must be  stable both  thermochemically and mechanically,  even when the volatility (41  diagram2,21,24,25  predicts  high  vapor  pressures  atm 5  101 325 Pa) in certain regions of the scale. We think the experiments15,16,21,22 suggest that with a porous ZrO2(c) skeleton, the high gas volatility problem can be avoided simply by the liquid phase retreating somewhat outwards,21 while still maintaining  mechanical integrity by holding onto the outer ZrO2(c) skeleton with capillary forces. The porous ZrO2(c) skeleton acts as a condensing substrate and mechanical support to the borosilicate  liquid. This  is not available when oxidizing monolithic SiC,  where the SiO2 (l) coalescence and layer  is mechanically unstable against gas bubble  shear-off or  spallation above  the SiO  boiling temperature.  In Section III, we make  some molecular-level predictions  regarding likely oxygen transport mechanisms  in the borosili cate liquid, based on density functional  theory (DFT) calcula tions. While  it  is  commonly  accepted  that molecular O2 low temperatures,26  permeate  through the glassy network at  the DFT  calculation  results  suggest that network incorp orated defects such as peroxyl linkage27,28 or oxygen-deficient centers29,30 will overtake molecular O2 as dominant oxygen carriers at above B15001C.  Y. Blum—contributing editor  This work was performed mostly at  the Ohio State University and supported by the  Ceramics  and Non-Metallic Materials Program in the Air Force Office  of  Scientific  Research (FA9550-05-1-0026).  Presented at  the AFOSR Workshop on Ultra-High-Temperature Ceramic Materials  hosted by SRI International, July 23-25, 2007.  w  Author to whom correspondence should be addressed. e-mail  liju99@alum.mit.edu  Manuscript No. 23766. Received September 19, 2007; approved December 12, 2007.  Journal  J. Am. Ceram. Soc., 91 [5] 1475 - 1480 (2008)  DOI: 10.1111/j.1551-2916.2008.02319.x  r 2008 The American Ceramic Society  1475  \\x0c', 'II.  Thermochemical and Mechanical Stability Analysis of the  Oxide Scale  ZrB21SiC has a complex scale structure containing at least two condensed phases: ZrO2(c), which in this section is assumed to be a highly porous skeleton with percolating holes, and a silica rich liquid phase  that wets  the  skeleton. Gas  species of  the  greatest  interest are B2O3, SiO, and CO, although BO, B2O2, B2O, etc. are also present and can be the dominant gas carriers in reducing conditions, and can play important roles in mass transport.31 SiO could evolve by for instance:  2SiO2 ðl Þ ! 2SiOðgÞ þ O2  (3)  (3) is a key reaction that has been used in constructing vola tility diagrams.2,21,24,25 When in contact with SiO2(l), with decreasing oxygen chemical potential or the equivalent partial pressure (pO2k), SiO will have higher equilibrium vapor pres(pSiOm). Volatility diagram of sure the ZrB21SiC system2,21 indicates that when T4TB \\x19 17751C, the peak equilibrium vapor pressure of SiO inside the scale could exceed 1 atm, which  would then induce a boiling transition (gas bubbles can nucleate  and grow inside the liquid). This violently disrupts the SiO2(l) scale in the case of oxidizing monolithic SiC. However, the scales  of ZrB21SiC and HfB21SiC appear to be much more tolerant of such a boiling transition. It is precisely in this T4TB regime that ZrB21SiC and HfB21SiC demonstrate oxidation resistance superior to monolithic SiC, which otherwise is considered a  highly oxidation-resistant material (see Fig. 1).  We  hereby  suggest  a  ‘‘dynamic  view’’  (Fig.  2(a))  and  a  ‘‘steady-state view’’ (Fig. 2(b)) of why ZrB21SiC is superior to monolithic SiC. The two views are inherently consistent. Imag ine a ZrB21SiC specimen is gradually being heated up in an oxygen-rich environment like normal air (pO2 5 0.2 atm), initially from T \\x1c TB. At such high ambient pO2, a protective SiO2(l) film will condense on top from the very beginning,24 that wets the ZrO2(c) skeleton, leaving no voids at the base. At ToTB, monolithic SiC in fact resist oxidation better than ZrB21SiC. Also at ToTB, the pSiO branch of the volatility diagram2,21 in contact with SiO2(l) has a formal thermodynamic definition but is not physically realizable as pure SiO gas bub bles, because any SiO gas bubble anywhere will be crushed by  the dual  forces of  surface  tension and hydrostatic pressure,  which we take to be 1 atm inside SiO2(l). However, as the temperature is brought up to T4TB, a sharp transition happens inside SiO2(l). Now SiO gas bubbles can nucleate at the base, with Dp \\x11 pSiO\\x001 atm40 working against the surface tension. The dynamic view (Fig. 2(a)) examines how the gas bubbles  grow and coalesce, paying attention to the role of  the ZrO2(c)  skeleton.  It is likely that the ZrO2(c) skeleton will regulate gas bubble dynamics. Unlike unconstrained growth inside a completely  liquid scale, gas bubbles (SiO, CO, B2O3, BO, O2, etc. mixture) in a semisolid porous scale are forced to grow into long fingers  (a pressure difference of  the order atm is large enough to dis place a liquid, but usually not enough to displace a solid). Re action (3) could then happen on one end of the gas finger, BO,  B2O3, SiO, CO, etc. would then diffuse along the gas finger with O2 diffusing in the opposite direction, and finally when pO2 gets high enough, SiO could get reoxidized to form SiO2(l) on the other end of the finger,22 and B2O3, etc. would get solvated in the liquid and continue to diffuse up the scale. This is equivalent  to a channeling transfer of SiO2(l) from one end of the finger to the other, which is mechanically untenable without the support  and constraint of the ZrO2(c) skeleton. In reality the pores are tortuous instead of straight, giving the effusing SiO(g)  much opportunity to react with O2(g) near SiO2(l) product collected on the ZrO2(c) substrate. The skeleton may also impart significant mechanical integrity to the  the  end, and the  scale  in the  case of bubble outbreaks  (bubble diameter  con strained by the pore diameter) or under external shear flow, be cause  SiO2(l) porous ZrO2(c) skeleton by capillary forces across a large contact area. In short,  adheres  strongly  to  a  highly  the  dynamical  view is  that  the ZrO2(c) silica-rich liquid, playing an important  skeleton  helps  to  collect and retain the  mechanical role.  There is also a ‘‘steady-state’’ explanation (Fig. 2(b)), beginning with the interpretation of volatility diagrams.24,25 Volatility  diagrams  represent  chemical  equilibria when assuming  some  volatile gas species are in contact with certain condensed phas es—solid or liquid. Solid or liquid makes a difference here, be cause a liquid could flow in or flow out, easily retreating from a  region if necessary. If a certain condensed phase retreats, then an  originally high volatile gas pressure—assuming the condensed  phase was there—loses its significance. For instance, pB2O3 in Fig. 11 of Opeka et al.2 must appear more important than it  really is in ZrB21SiC oxidation, because we know there is no B2O3(l) to make contact with in reality at these temperatures. Another way of seeing this is that in reality all condensed-phase  B2O3, if they exist, are solvated inside SiO2(l), with much lower activity than in pure B2O3(l), and thus the actual pB2O3 vapor pressure should be lowered in proportion and will not be as  dangerously high as it originally looks. This understanding of  the volatility diagram may be translated into the following rule:  liquid phases will retreat from the region where some volatile gas  species, were  they  in  contact, will  have  high  vapor  pressure  (approaching hydrostatic pressure  inside  the  liquid),  to region  of lower vapor pressure according to the volatility diagram; after  the retreat, the actual vapor pressure of the volatile gas species will  200 150 100 80 60 40 30 20 15  10 8 6 4 3  2 1.5  2050  2100  2150  2200  2250  2300  2350  2400  2450  Temperature,°K  d  (  l  g o  c s  a  l  e  )  i  n  m  l i  s  (  1  h  r  .  )  HfB2  HfB2 − 20 % SiC   SiC  Fig. 1.  One-hour oxidation resistance of monolithic HfB2, monolithic SiC, and HfB21SiC composite (taken from Clougherty, Pober, and Kaufman).7  Fig. 2.  The  ‘‘dynamic view’’  (a) and ‘‘steady-state view’’  (b) of why  ZrB21SiC has better oxidation resistance than monolithic SiC above the SiO boiling transition temperature.  1476  Journal of the American Ceramic Society—Li et al.  Vol. 91, No. 5              \\x0c', 'be automatically lowered than what  the  volatility diagram has  originally indicated for the evacuated region.  The above rule comes from thermochemistry. Applying the  rule to oxidizing monolithic SiC above TB (see Fig. 2(b) middle), we see that the only thermochemically sound steady-state ar rangement  is  for SiO2(l) to the low pSiO-high pO2 region, and let gas-phase diffusion take over in the intervening  to retreat  gap, where there will be no condensed liquid phase,  thus shut ting down the high gas volatility. Unfortunately, although this  setup is thermochemically and diffusion-kinetically sound,  it is  obviously mechanically unstable. The ‘‘scale’’ would easily shear  off or  spall. This is fundamentally because in monolithic SiC,  with  only  a  single  condensed-phase  oxide  product, which  is a liquid, there is no way to satisfy both thermochemical and mechanical stabilities simultaneously at T4TB. In contrast, when oxidizing ZrB21SiC, one gets densed-phase oxide products. ZrO2(c) itself has only low volatilities of ZrO(g) and ZrO2(g) (see Fig. 11 of Opeka et al.2). Furthermore it is a solid. So it does not retreat  two con from the high  SiO volatility region, maintaining mechanical connection with  the main body. The silica-rich borosilicate liquid duly retreats  from the high SiO volatility region, thereby removing the high  SiO volatility automatically. This staggered placement of solid and liquid phases at T4TB, with internal gas finger diffusion and gas-solid reactions (Fig. 2(b) bottom), is both thermochem ically and mechanically sound if  the ZrO2(c) enough, such that even after the retreat the liquid phase can still  skeleton is  long  hold onto the solid by capillary forces, to have a ‘‘solid pillars,  liquid roof’’ architecture as suggested by many experiments.15,16,21,22 The actual oxidation of SiC particles at T4TB then no longer follows reaction (2), but directly  SiCðcÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  (4)  a gas-solid reaction without going through the liquid phase,22  which will lead to a porous ‘‘SiC-depleted’’ substrate layer in the ZrB21SiC.13 The ZrO2(c) ‘‘solid pillars’’ will also grow longer at the base by gas-solid reaction, with O2(g) reactant and BXOY(g) product. On the other end of the gas finger, we will have the  reverse of reaction (3):  2SiOðgÞ þ O2 ! 2SiO2 ðl Þ  (5)  which replenishes the ‘‘liquid roof ’’.  The above explanation,  if correct, could rationalize why the  microstructure of ZrB21SiC is important for its oxidation resistance,20 because the pore sizes of ZrO2(c) might be related to the preoxidation SiC particle sizes. As Gasch et al. mentioned,13  ‘‘at 20 volume percent SiC, if the SiC particles are assumed to be  small spheres randomly distributed throughout the HfB2 matrix, the amount of SiC should be above the percolation threshold.  This means that the SiC particles form a network that is inter connected in three dimensions.’’ For  a  certain fixed volume  fraction,  smaller pores and better  connectivity inside ZrO2(c) could enhance the collection and retention of the silica-rich liq uid. Thus, nanoscale SiC particles might improve the oxidation resistance of ZrB21SiC,20 by refining the microstructure of the in situ formed ZrO2(c) skeleton. The proposed view also explains why 70-80 vol% of the com posite is dedicated to ZrB2. It is seen that a porous ZrO2 could be advantageous for the oxidation resistance for a mechanical reason  (not a diffusion kinetics one), if there is also a liquid oxide prod uct phase to ‘‘collaborate’’ with. The porous skeleton needs to be  strong enough as well as sufficiently long, to have enough room  for the liquid oxide phase to retreat outward. Otherwise, the two  phases may still have to separate (‘‘high volatility blows away the  liquid roof’’), and the entire system would lose oxidation protec tion. One cannot have too much SiC (and thus the liquid) and not  enough oxide skeleton, and maintain the mechanical and the rmochemical  stabilities of  the ‘‘solid pillars,  liquid roof’’ archi tecture. Because capillary force holds the solid and liquid phases  together, microstructural refinement of the oxide scale will lead to  stronger capillary adhesion per unit volume, which could lead to  significant improvement of the overall oxidation resistance.  III.  Oxygen Transport in Silica-Rich Liquid  From Section II model, we see that if ZrO2(c) has a percolatingholes microstructure, the borosilicate liquid will define the effec tive barrier against oxygen,  irrespective of whether fully dense  ZrO2 against oxygen or not than the borosilicate liquid, as the liquid  is  intrinsically better barrier  (in open-circuit  condition)  occupies both serial and parallel oxygen transport pathways in  the ‘‘solid pillars,  liquid roof’’ architecture.  In this  section we  focus on the atomic-level events that govern oxygen transport in  the borosilicate liquid.  Experimentally,  it is still challenging to accurately determine  the composition profile of the borosilicate liquid because boron  is a light element. According to the Hertz-Knudsen-Langmuir equation,32 the net evaporation flux of a species from a liquid  surface is  J ¼ aDp=ð2pmkBT Þ1=2  (6)  Where Dp is the difference between the equilibrium vapor pres sure and the actual vapor pressure of the species at the surface, and a is a coefficient of order 1. Because pure B2O3(l) has much higher equilibrium vapor pressure than pure SiO2(l) in an oxygen-rich environment, a borosilicate liquid facing air would  preferentially evaporate B2O3 instead of SiO2. Thus the borosilicate liquid should be overall silica rich, with a composition  gradient  that  is B2O3 depleted at B2O3 enriched at the gas finger-liquid interface (Fig. 2(b)), as B2O3(g) and other boron-bearing gas species31 are carried along with SiO(g) in the gas finger and get absorbed into the liquid.  the liquid-air  interface, and  Previously, Bongiorno and Pasquarello have studied oxygen transport in pure silica glass28,33,34 using a multiscale modeling  approach that combines high-level quantum mechanical (DFT)  calculations of  the diffusing oxygen species and local  energy  barriers, with  kinetic Monte Carlo  sampling  of  connected  migration pathways. To bound the results, we decide to model  a  borosilicate  liquid  composition  of  equal B2O3  and  SiO2  proportions.  One major challenge in modeling any glass or liquid is to have  reliable atomic structures. We have adopted the structure generation approach of Van Ginhoven, Jonsson, and Corrales,35  which was  shown to reproduce experimental pair distribution  functions for pure silica. The approach requires a classical  in teratomic potential  to perform long-time molecular dynamics  (MD)  simulations at  the beginning,  followed by further DFT  optimizations. To generate the classical potential, we adopt the van Beest, Kramer, and van Santen parameterization36 for Si-O  interactions, but fit B-O and Si-B interactions  to a series of  small DFT calculations for bulk B2O3, using the software package GULP.37 Then, starting from random positions for oxygen,  boron and silicon atoms in the supercell, we perform a sequence  of classical MD simulations at  temperatures 6000, 5000, 4000,  3000, 2000, and 1000 K. The resulting structures were then used  as input geometries for further DFT calculations (Vienna ab initio simulation program38,39 with spin-polarized PW91 functional,40 projector augmented wave method,41,42 planewave ki netic  energy cutoff 400 eV). For our  initial  studies, a cubic  supercell containing 14SiO217B2O3 formula units is used, with total 77 atoms. The average density is 2.3 g/cm3. A typical liquid structure at T 5 25001C, after further equilibration by ab initio  MD,  is shown in Fig. 3(a). It clearly has a framework structure  with no long-range order, and contains with no dangling bonds  (all Si are fourfold coordinated to O, and all B are threefold  coordinated to O).  Based on these atomic structures, we have studied the ther modynamic stability and diffusion kinetics of solvated oxygen \\x03 and atomic O\\x03 in borosilicate liquid. The former molecules O2 stay inside the open cages of the framework and do not interact  May 2008  Thermochemical and Mechanical Stabilities of the Oxide Scale of ZrB21SiC  1477  \\x0c', 'chemically with the framework. The latter are chemically incor porated into the network Si-O-O-B27,28 (Fig. 3(b)  in  the  form of  peroxyl  linkage  inset), extra bridging O between two  B (Fig. 3(b) inset), and others. Because diffusion could be a rare  event,  simply performing MD simulations  and tracking  the  mean squared displacements may not be sufficient, and ener gy-landscape exploration techniques such as nudged elastic band (NEB)43 calculations may be needed. These methods com pute the minimum energy path (MEP) and saddle-point config uration of thermally activated processes, transition-state theory44 to estimate the rates. \\x03 inside a cage and perform DFT MD First, we place one O2 simulation at 25001C for 11 ps (see movie S1 at http://  and  then  use  alum.mit.edu/www/liju99/Papers/08/JACerS/).  It  is very  clear  from the MD trajectory (in contrast  to those of simple liquids  such as Ar, as well as water) that the borosilicate liquid still maintains a very ‘‘rigid’’ framework at 25001C. Indeed, pure silica is a strong glass-forming liquid,45 and its viscosity does not  show a precipitous drop above the glass transition temperature. For instance, even at 25001C, pure silica still has a shear viscosity ZB104 poise, which is a million times thicker than that of  room-temperature water. This means the liquid still has a well defined network structure at any given moment, and its topo logical change does not occur at  the same timescale as,  for in stance, its own Si-O-Si bond stretching. Also, from the 11 ps ab \\x03 is essentially trapped inside initio MD trajectory, we find the O2 one cage. It just bounces back and forth many times inside a  jiggling cage, with no possibility of escape within the MD sim ulation timescale. These facts suggest that one is still justified to  use transition-state theory and numerical schemes like the NEB \\x03, despite it is embedded method43 to characterize diffusion of O2 in a liquid. From our MD simulations, adding 50% B2O3 to silica does not seem to change this consideration qualitatively.  A complication for  the NEB calculation is  that unlike  in  crystals, diffusion inside an amorphous framework has a distribution of local minima and activation energies.33,34 Not all cages  have the same volume, nor the same energy for opening up the \\x03 squeezes from one cage constrictions (see movie S2) when O2 into the other. In Fig. 5 of Bongiorno et al.,33 we have shown a \\x03 diffusion (vehicular typical DFT-NEB-calculated MEP of O2 diffusion mode). The forward hop barrier is 1.8 eV, whereas the  backward hop barrier is 1.4 eV. These are somewhat higher than  the 1.12 eV effective migration barrier that Bongiorno and Pas\\x03 vehicular diffusion in pure silica, quarello34 predicted for O2 perhaps because of the B2O3 modifications to the network. More calculations are needed in order to have better statistics.  For vehicular diffusion inside a liquid, there is a well-known  Stokes-Einstein relation:  DðOn 2 Þ \\x18 kBT =6pZRðOn  2 Þ  (7)  \\x03) is a nominal hydrodynamic radius of the molecule. where R(O2 Even though the Stokes-Einstein relation is quite successful in simple liquids, it can fail in network-forming liquids.46,47 Norton  measured the permeation of gaseous oxygen through vitreous silica and found an activation energy of 27 kcal/mol (1.17 eV)48 \\x03), for D(O2 in substantial agreement with later measurements.26 However, the activation energy governing Z, the viscosity of is in the range of 5.3-7.5 eV.49 So clearly Eq. (7)  vitreous silica,  does not work. From our DFT modeling, the physics governing \\x03 vehicular diffusion is seen to be an the activation energy of O2 elastic deformation of the framework (elastic opening of the  constrictions,  see movie S2) without changing its network the physics behind the activation energy of Z  topology. But  must  involve network topology changes, which necessarily in volve Si-O bond breaking. \\x03 vehicular diffusion, oxygen transport may In addition to O2 also occur by Grotthuss-type oxygen-hopping mechanisms,50 linkage27,28 or oxmediated by network defects such as peroxyl centers.29,30 These mechanisms would  ygen-deficient  involve  bond breaking, and typically higher effective activation ener gies—mostly due  to the  formation energies of  such defects. \\x03 breaks up Figure 3(b) shows one such pathway, where an O2 into two O\\x03:  2 ! On þ On On  (8)  The two O\\x03s then move independently of each other for a while, and eventually recombine on the other side of the cage (see  movie S3) in this particular NEB calculation setup. In node 3 of the calculated MEP (Fig. 3(b) inset), one O\\x03 takes the form of a peroxyl linkage Si-O-O-B, while the other O\\x03 takes the form of an additional B-O-B bridge, making the two boron atoms four fold coordinated.  We find from multiple NEB calculations in the borosilicate  framework that  these coordination defects, once formed, can  interconvert easily, suggesting low migration barriers, similar to  0  1  2  3 4 5 NEB node number  6  7  8  0  0.5  1  1.5  2  2.5  3  3.5  4  4.5  5  (a)  (b)  E  n  e  r  y g  [  e  V  ]  Dissociation-Recombination pathway:  O * O* + O* O * :  oxygen: red boron: orange silicon: silver  Fig. 3.  (a) A typical borosilicate liquid structure (14SiO217B2O3, plus \\x03) generated from a sequence of long-time classical moone solvated O2 lecular dynamics simulations, followed by density functional theory (DFT) molecular dynamics simulation at 25001C. Note that the super cell has been replicated three times in all  three dimensions to facilitate  visualization of the framework. (b) A typical DFT calculated minimum  energy path (MEP) of dissociation-migration-recombination reaction \\x03, where oxygen exthe borosilicate framework: O2 \\x03 and the network occurs. The inset shows the atomic change between O2 configuration (replicated for visualization) of node 3 on the MEP, which  inside  \\x03-2O\\x03-O2  contains one peroxyl  linkage (Si-O-O-B) and one extra B-O-B bridge.  Oxygen: red, boron: orange, silicon: silver.  1478  Journal of the American Ceramic Society—Li et al.  Vol. 91, No. 5    \\x0c', 'interstitial  defects  in metals. On  average,  the  right-hand  side (RHS) of reaction (8)  is about 2.8 eV higher in potential  energy than the left-hand side (LHS), as indicated by node 3, 4,  5, 6 of Fig. 3(b) MEP (see also movie S3). However, there are the two O\\x03s, once formed, can move independently of each other inside the liquid.  two free translational centers on the RHS:  The LHS  has  only  one  free  translational  center,  in  order  to maintain its molecular form. Thus the RHS of reaction (8)  has  entropic  advantage  of  approximately  kBlnc, concentration (per  where  c  is  the  prevalent  oxygen  carrier  formula  unit of SiO21B2O3), whereas The classic enthalpy-entropy  the LHS has energy advantage.  tradeoff  in  free  energy  then  suggests  that  there  exists a temperature TC, below which O2 O\\x03s concentration, and above which in an oxygen-rich environ \\x03  is  dominant  in  are dominant,  if  the borosilicate is  ment (the equivalent pO2 is high). In lower equivalent pO2 environment, oxygen-deficient centers29,30 are also possible carriers.  In reference to O2 in the gas phase, we find the average po\\x03 solvated in borosilicate is 0.73 eV per moltential energy of O2 ecule, which is essentially the elastic energy in the framework  needed to accommodate the molecule. The average potential energy of O\\x03, on the other hand, is about 1.78 eV per O. These \\x03 and O\\x03 in energies are used to compute the concentration of O2 borosilicate in equilibration with pO2 5 0.1 atm, shown in Fig. 4. Note that scales linearly with pO2, whereas scales linearly with (pO2)1/2. Thus, deeper and deeper into the scale, as the oxygen chemical potential gets lower and lower,  cO\\x03  2  cO\\x03  the Grotthuss-type oxygen transport  should become relatively  more and more important.  \\x03 in pure silica at 10781C and pO2 5 1 atm to be 1.9 \\x02 10 Norton measured the solubility of O2 cm3 STP O2 gas per cm3 silica.48 This amounts to 5.1 \\x02 1016 per cm3 silica or a dimensionless concentration of about c 5 10\\x005, at pO2 5 1 atm. This is about two orders of magnitude higher than our DFT-predicted O2 \\x03 solubility in borosilicate, shown in Fig. 4. This could be due to the structural difference between pure silica framework and borosilicate framework. As Bongiorno and Pasquarello34 noted, \\x03 depends sensitively on the cage inthe potential energy of O2 terstice volume, and there is certainly a structural difference be \\x003  tween pure silica and B2O3-modified networks. This could also be partly due to the intrinsic errors of PW91 density functional,40 which are known to give large errors treating isolated mol ecules  (the reference state), and nonbonding interactions  (O2  \\x03  interactions with the framework).  The above uncertainties aside, it is still a rather conservative estimate that above TCB15001C the dominant oxygen carriers in the borosilicate ‘‘liquid roof’’ are network defects instead of \\x03, because molecular O2  there are other  factors not  shown in  Fig. 4 that disfavors the Deal-Grove mechanism: (a) the DFT calculations indicate that the O\\x03s not only have entropy advantage, but also mobility advantages over O2 (b) intense aerothermal heating environment may introduce a significant level of dissociated oxygen on the outer liquid surface,14,33 which would favor O\\x03 diffusion from a nonequilibrium kinetics perspective, (c) as pO2 drops from B104 Pa on the outer liquid surface to o 10\\x005 Pa equivalent at the internal gas finger-liquid interface (Fig. 2(b)) according to the volatility diagram,2,21 the balance \\x03 vehicular diffusion to will shift more and more away from O2 network defect Grotthuss diffusion. Oxygen-deficient centers29,30 may play a significant role in oxygen transport at  \\x03,  low  equivalent pO2s. It seems plausible that oxygen in the borosilicate liquid could react with the underlying substrate or SiO(g),  injecting oxygen vacancies (e.g., regions of high B and Si stoic hiometry) into the liquid, which then diffuse up the scale to re\\x03 or O\\x03 somewhere inside the liquid. (d) In all combine with O2 the DFT calculations, we have only considered neutral network  defects. Charging the defects may significantly greatly reduce their formation energies,30,51 although we will then need to solve  the complementary problem of what other defects compensate  the charge and carry out ambipolar diffusion under open-circuit  condition.  \\x03 suggests The mobility advantage of the peroxyl linkage over O2 that arc-jet testing,13,33 which introduces a nonequilibrium distri bution of dissociated oxygen atoms on the surface, will  likely lead  to faster oxidation than ordinary furnace testing at the same tem perature. Also, because the network defects are chemically incor porated into the network and thus interact more strongly with \\x03, small changes in the glass chemistry could lead to solutes than O2 large changes in the oxygen diffusivity by defect trapping/gettering,30 much more than what Eq. (7) could have suggested.  IV.  Summary  We present a congruent explanation of the oxidation protection  of ZrB21SiC based on a ‘‘solid pillars, tecture, where the borosilicate liquid defines the effective diffu liquid roof’’ scale archi sion barrier, and the solid zirconia collects and retains the liquid  and provides mechanical support. Internal gas fingers will form  as  the liquid phase retreats  to remove the high SiO volatility  above a boiling transition temperature. At such high tempera tures, to satisfy both thermochemical and mechanical stabilities,  the ‘‘solid pillars,  liquid roof’’ architecture seems to be a viable  solution, not available to monolithic SiC.  Compared  with  the  borosilicate  liquid  phase,  whether  fully dense  zirconia  is blocking or unblocking  to oxygen in  open-circuit  condition depends on its  electronic  conductivity  (transference number), which in turn depends on how the charge  defects are compensated inside the crystal, related to the amount  of impurities. If the zirconia phase has a highly porous micro structure, however, then the above discussion is likely irrelevant  and the borosilicate liquid phase will control the effective diffu sion barrier, because it will occupy both serial and parallel ox ygen transport pathways.  At low temperatures, it is commonly accepted that molecular \\x03 dominates oxygen transport. However, oxygen O2 principles calculations with detailed borosilicate atomic structures,  from first it seems unlikely that this will remain the case at temperatures of  practical interest for the ZrB21SiC thermal protection system (above 15001C). This means that the oxidation rate will likely  have a complex, sublinear dependence with respect to the external  oxygen partial pressure. Also, if the oxygen carriers are chemically  incorporated and interact strongly with the framework,  there is  hope that by tuning the glass composition, the carriers could be  trapped, thereby slowing down oxygen diffusion.  Acknowledgments  We would like to thank Mark Opeka, Inna Talmy, Robert Rapp, Triplicane  Parthasarathy, Ronald Kerans, and Nitin Padture for instructive discussions.  1000  1500  2000  2500  10 -10  10 -9  10 -8  10 -7  10 - 6  T [K]  C  n o  e c  n  t  r  a  i t  n o  pO2 = 0.1 atm  c  Fig. 4.  \\x03 and O\\x03 (per formula unit of Estimated concentrations of O2 SiO21B2O3) inside borosilicate framework, in equilibration with pO2 5 0.1 atm.  May 2008  Thermochemical and Mechanical Stabilities of the Oxide Scale of ZrB21SiC  1479  \\x0c', 'References  1W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. Zaykoski,  ‘‘Refrac tory Diborides of Zirconium and Hafnium,’’ J. Am. Ceram. Soc., 90 [5] 1347-64  (2007). 2M. M. Opeka, I. G. Talmy, and J. A. Zaykoski,  ‘‘Oxidation-Based Materials  Selection for 2000 Degrees C Plus Hypersonic Aerosurfaces: Theoretical Consid erations and Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 3A. K. Kuriakose and J. L. 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},{
  "_id": 271,
  "PDF": "Thermodynamic Analysis of ZrB2–SiC Oxidation Formation of a.pdf",
  "Text": "['Journal  J. Am. Ceram. Soc., 90 [1] 143 - 148 (2007)  DOI: 10.1111/j.1551-2916.2006.01329.x  r 2006 The American Ceramic Society  Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a  Department of Materials Science and Engineering, University of Missouri-Rolla, Rolla, Missouri  SiC-Depleted Region  William G. Fahrenholtz*,w  A thermodynamic model was developed to explain the formation  of a SiC-depleted layer during ZrB2-SiC oxidation in air at 15001C. The proposed model suggests that a structure consisting of (1) a silica-rich layer, (2) a Zr-rich oxidized layer, and (3) a  SiC-depleted  zirconium diboride  layer  is  thermodynamically  stable. The SiC-depleted layer developed due to active oxida tion of SiC. The oxygen partial pressure in the SiC-depleted layer was calculated to lie between 4.0 \\x02 10\\x0014 and 1.8 \\x02 10\\x0011 Pa. Even though SiC underwent active oxidation, the overall  process was consistent with passive oxidation and the formation  of a protective surface layer.  I.  Introduction  ULTRA-HIGH-temperature ceramics (UHTCs) have been proposed as candidates for applications such as thermal protection materials on hypersonic aerospace vehicles and re-usable atmospheric re-entry vehicles.1,2 Among the UHTCs, both zir conium and hafnium diborides (ZrB2 and HfB2) have a desirable combination of high melting temperature (430001C), resistance  to chemical attack, and other physical properties that make them attractive for these applications.3 Historically, the process ing, properties, and oxidation behavior of UHTCs were inves tigated in Russia and the U.S. during the 1950s, 1960s, and 1970s.4,5 The recent drive to develop hypersonic flight vehicles  has led to a resurgence of interest in UHTCs. Groups in Italy,  the U.S., and Japan are currently investigating microstructure- systems.6-12  processing-property  relations  UHTC  For  in  thermal protection applications on hypersonic flight vehicles, UHTCs will be exposed to high temperatures (15001C and above) and oxidizing environments.13,14 This paper focuses on  ZrB2-based ceramics for these applications, but HfB2-based ceramics exhibit analogous behavior and could be described with  the same type of model. (cr)z Exposure of ZrB2 oxidation to ZrO2 (cr) and B2O3 about 11001C, the B2O3 (l) forms a continuous temperature regime, parabolic (diffusion controlled) kinetics are  (l) by Reaction (1). Below  stoichiometric  In this  results  in the  to air  layer.  observed with reported activation energies in the range of 80- 120 kJ/mole,15,16 which are consistent with reported values for oxygen diffusion in B2O3 (l).17 Based on the activation energy and the dependence of oxidation rate on the partial pressure of oxygen (pO2 ), the rate of oxidation below 11001C appears to be controlled by the transport of oxygen through B2O3 (l).16,18-20 The ZrO2 appears to form a porous skeleton that does not enhance the oxidation protection, but may provide mechanical  N. Jacobson—contributing editor  Manuscript No. 21450. Received February 2, 2006; approved August 2, 2006.  This work was supported by the Ceramics program in the Division of Materials Re search at the National Science Foundation (DMR-0346800) and by the Ceramic and Non Metallic Materials Program in the Air Force Office of Scientific Research (FA9550-06-1 0125).  *Member, American Ceramic Society.  w  z  Author to whom correspondence should be addressed. e-mail: billf@umr.edu  The NIST-JANAF convention is used in this paper whereby (cr) indicates crystalline  solids, (l) indicates liquids or amorphous solids, and (g) indicates gaseous species.  143  integrity to the liquid B2O3 scale.  ZrB2 ðcrÞþ5  2O2  ðgÞ ! ZrO2 ðcrÞ þ B2O3  ðl Þ  (1)  Above 11001C, the oxidation rate increases compared with what  would be predicted from the diffusion-controlled (parabolic) behavior observed at lower temperatures.19 Between 11001 and 14001C, so-called para-linear kinetics are observed in which the  overall rate of mass change is a combination of weight gain due  loss due to  to formation of B2O3 (l) and ZrO2 (cr) and weight volatilization of B2O3 (l).19,21 Above 14001C, the rate of evaporation of B2O3 (l) is greater than its rate of production, leaving a non-protective porous ZrO2 (cr) scale. At these temperatures, ZrB2 (cr) oxidation in air exhibits rapid linear kinetics.22,23 Despite the evaporation of B2O3, oxidation of ZrB2 (cr) above 11001C results in mass gain as the mass of ZrO2 greater than the mass of ZrB2 consumed. At 11001C and above, the oxidation resistance of ZrB2 (cr) can be improved by adding SiC (cr) to promote the formation of a silica-rich scale.24-27 Below B11001C, SiC (cr) additions do  formed is  not affect the oxidation rate or the composition of the protective layer.27 The oxidation of SiC (cr)  is much slower than that of  ZrB2 (cr) in this temperature regime, as evidenced by SiC (cr) inclusions present in the scale.19 Above B11001C, SiC (cr) oxidizes by Reaction (2) to form SiO2.19 The silica-containing scale on ZrB2-SiC is stable at higher temperatures than the B2O3 scale on ZrB2 because of the lower volatility of SiO2 (l) compared with B2O3 (l) at these temperatures. Thus, ZrB2-SiC exhibits slow, diffusion-controlled mass gain kinetics over a much greater temperature range than has been reported for pure ZrB2.24  a-SiC ðcrÞþ3  2O2  ðgÞ ! SiO2  ðl Þ þ CO ðgÞ  (2)  The formation of a ‘‘SiC-depleted layer’’ during oxidation of  ZrB2-SiC and HfB2-SiC has been noted by several authors.23,24,26,28 This layer forms below the outer silica-rich scale  and it consists of unoxidized diboride, but may contain some  retained SiC (cr). As shown in Fig. 1, oxidation of ZrB2-SiC at 15001C produces a surface scale that is SiO2 rich compared with the scale on pure ZrB2. Under the conditions of this experiment (heating at 101C/min to 15001C in air, hold 30 min), the under lying layer was composed of unoxidized ZrB2 (cr). Under some conditions, an additional layer consisting of crystalline ZrO2 (cr) and/or amorphous SiO2 (l) has been observed between the SiCdepleted zone and the outer glassy scale.8 It has been reported  that this layer also contains boron (B), but the form or phase in which it exists was not specified.23,26 The distribution of B was  not quantified in the present analysis due to the low sensitivity of  energy-dispersive spectroscopy to light elements.  Layers depleted of SiC (cr) have also been observed under the  more severe conditions imposed by arc heater testing for both ZrB2-SiC and HfB2-SiC.10,23 Both ZrB2-SiC and HfB2-SiC continue to exhibit protective behavior in these extreme envir onments. Apparently, SiC (cr) is removed from the ZrB2 matrix as gaseous species (e.g., Reaction (3)) as neither SiO2 nor other condensed phases have been observed in the SiC-depleted layer.  However,  to date,  the  thermodynamic  conditions  that  pro mote  the  formation of  the SiC depleted layer have not been  \\x0c', '144  Journal of the American Ceramic Society—Fahrenholtz  Vol. 90, No. 1  Fig. 1. A cross-sectional image (a) and elemental maps showing the distribution of (b) O, (c) Si, and (d) Zr indicating the formation of a surface layer containing Si and O and a SiC-depleted layer after heating zirconium diboride (ZrB2)-30% SiC to 15001C for 30 min in air. The silica layer and the SiCdepleted layer are both B10 mm thick.  thoroughly explored.  SiC ðcrÞ þ O2 ðgÞ ! SiO ðgÞ þ CO ðgÞ  (3)  The purpose of this paper is to provide thermodynamic jus tification for the formation of a SiC-depleted layer during oxi dation of ZrB2-SiC ceramics. A thermodynamic model consistent with reported oxidation behavior is developed and  that  is  discussed.  II.  ZrB2 Volatility Diagram at 15001C  Volatility diagrams are a concise means for understanding gas-solid interactions such as oxidation reactions.29 Essentially,  volatility diagrams are plots of  the vapor pressure of  the pre and temperature.  dominant gaseous species as a function of pO2 Volatility diagrams have been prepared for materials such as SiO2, SiC, Si3N4, MgO, and MgO-C.30,31 In addition, volatility diagrams have been used to understand the stability of oxide scales in ultra-high-temperature systems.1 More recently, a ZrB2 volatility diagram was compiled.22 By comparing the ZrB2 volatility diagram with the SiC volatility diagram, it may be possible  to understand the formation of  the SiC-depleted layer from a  thermodynamic point of view.  The construction of a ZrB2 volatility diagram has been described previously.22 The same methodology was used to compile the 15001C volatility diagram for ZrB2 for this paper (Fig. 2) using data from the NIST-JANAF tables.32 It should be noted  that  this diagram was calculated assuming that no water was  present. The stability of various HxByOz (g) species would most likely lead to increased attack in humid environments.33  The ZrB2 volatility diagram shows that the equilibrium pO2 for the ZrB2 to ZrO2 plus B2O3 transition (Reaction (1)) is 1.8 \\x02 10\\x0011 Pa (log pO2 5 \\x0010.74) at 15001C. The hibits four vapor phase transitions. Coincidentally,  system ex the vapor  transition from B2O2 (g) to B2O3 (g) at an oxygen partial pressure of 1.7 \\x02 10 \\x0011 Pa falls at almost exactly the same pO2 as (1.8 \\x02 10\\x0011 Pa). Consequently, Reaction (1) the vapor transition cannot be distinguished from the ZrB2 to ZrO2-B2O3 transition in Fig. 2. Note that the partial pressure of B2O3 (g) does not vary with pO2 as oxygen is neither consumed nor produced by the vaporization of B2O3 (l) to B2O3 (g). For exposure of ZrB2 to air at 15001C, the volatility diagram indicates that ZrO2 (cr) and B2O3 (l) form. Further, the partial pressure of the predominant gaseous species, B2O3 (g), is 209 Pa. Based on its vapor pressure, B2O3 (g) should evaporate at 15001C, leaving behind ZrO2 (cr). Below 11001C, ZrB2 exhibits parabolic mass gain kinetics consistent with diffusion of oxygen  Fig. 2.  Volatility diagram for zirconium diboride (ZrB2) at 15001C.  \\x0c', 'January 2007  Thermodynamic Analysis of ZrB2-SiC Oxidation  145  through the layer of B2O3 being the rate-limiting step as expected for the formation of a stable, protective surface oxide.16 However, at 15001C, ZrB2 which is consistent with the evaporation leaving a non-protective ZrO2 (cr) scale.4,15-21  linear oxidation kinetics,  exhibits  rapid,  B2O3  (cr)  of  the  III.  SiC Volatility Diagram at 15001C  The  thermodynamics of SiC oxidation have been thoroughly  explored in the literature and SiC volatility diagrams have been produced.30,34,35 A diagram based on the work of Heuer and  Lou was compiled for this paper using data from the NISTJANAF tables.33 Briefly, Heuer and Lou30 produced volatility diagrams by considering the stability of b-SiC (cr) under various  partial pressures of CO (g). For  the diagram in this paper  (Fig. 3), the partial pressures of CO (g) and SiO (g) in equilibrium with a-SiC (cr) were set equal based on the stoichiometry  of Reaction (3). At oxygen partial pressures greater  than the  SiC-SiO2 equilibrium, the presence of CO (g) was neglected. In addition, the SiO2 (l) formed by the reaction was assumed to be a pure material (i.e., aSiO2 5 1), even though some B2O3 (l) was undoubtedly present during ZrB2-SiC oxidation. The calculations in this paper also assumed that the SiC in ZrB2-SiC was in its alpha variant as a-SiC precursor powders were used to pre pare the experimental specimens. Thus, some lines in Fig. 3 are  shifted slightly compared with the Heuer and Lou diagram. For exposure of a-SiC (cr) to air at 15001C, Fig. 3 indicates  that SiO2 (l) should form. The predominant vapor species is SiO2 (g) with a partial pressure of B10\\x004 Pa. Under these conditions, SiO2 (l) should form a protective scale on a-SiC (cr) and pressure predicted for SiO2 (g), B10\\x004 Pa, and the lower partial parabolic mass gain kinetics would be expected. The low partial pressure predicted for SiO (g), B10\\x005 Pa, are consistent with formation of a stable, protective scale. Experimentally, SiC ex hibits passive oxidation with parabolic kinetics in air at 15001C.34 For ZrB2, the transition from a protective to a nonprotective scale occurred due to B2O3 (l) vaporization above 12001C.19 For a-SiC (cr), the transition from protective behavior  with parabolic mass gain kinetics to active oxidation with linear  mass loss kinetics occurs by a different mechanism. Above B16001C, the oxide scale on SiC is no longer protective in oxi dizing environments because the SiO2 (l) reacts with the underlying a-SiC (cr) to form SiO (g) and CO (g) by35  SiC ðcrÞ þ 2 SiO2  ðl Þ ! 3 SiO ðgÞ þ CO ðgÞ  (4)  IV.  Comparison of the ZrB2 and SiC Volatility Diagrams  The response of ZrB2-SiC (cr) predicted by comparing the  to oxidation at 15001C can be  respective  volatility  diagrams  (Fig. 4). Doing so assumes  that  the condensed phases do not  interact significantly and, thus, have unit activity. This assump tion is valid for the two base materials (ZrB2 and SiC) as they do not react or form significant solid solutions. In addition, it was  assumed that the oxide reaction products (ZrO2, SiO2, or B2O3) do not form compounds or solutions. Even though some phase diagrams suggest that ZrSiO4 is stable at 15001C,36 experimental oxidation studies for ZrB2-SiC do not report the formation of significant amounts of ZrSiO4 21,38; thus, ZrO2 activity was taken as unity. If SiO2 and B2O3 were to form a liquid solution, then the presence of B2O3 (l) would reduce the activity of SiO2 (l). The analysis in this paper assumed unit activity for SiO2 (l) the scale at 15001C because of the preferential evaporation of  in  B2O3 (l) due to its higher vapor pressure. Additional support for this assumption is offered by a recent report in which secondary  ion mass spectrometry (SIMS) detected a B content of less than  1 wt% in the outer SiO2-rich scale on ZrB2-SiC oxidized at 15001C in air.37 If compounds or solutions were to form, the  activity of  the compounds would be reduced, which would,  in  most cases, reduce the vapor pressures and shift the equilibrium  oxygen partial pressures  for Reactions  (1)  and (2)  to lower  values. Based on the assumptions that were made in the pres ent study, the analysis represents a limiting case, which can be  modified when more detailed information regarding the chem ical species is available. From Fig. 4, a-SiC (cr) oxidizes at lower oxygen partial pressures than ZrB2 (cr) as the line separating a-SiC (cr) from SiO2 (l) lies to the left of the line separating ZrB2 (cr) from ZrO2 (cr) and B2O3 (l). From the previous analyses, predicted oxygen partial pressures for oxidation of ZrB2 (cr) and a-SiC (cr) are B10\\x0011 and B10\\x0012 Pa, respectively. In addition to partial pressures for oxidation of ZrB2 and SiC, the overlapped volatility diagram provides information that can  be used to interpret experimental observations of the phases and  structures formed. From the combined diagram, exposure of ZrB2-SiC to air at 15001C is predicted to produce three condensed phases: ZrO2 (cr), B2O3 (l), and SiO2 (l). Experimentally, the scale on ZrB2-SiC resembles the scale on pure ZrB2 (cr) that consists of porous ZrO2 (cr) and a continuous amorphous indicating that, if B2O3 (l) geneous glass with SiO2 when cooled to room temperature. As discussed in the introduction, ZrB2 (cr) exhibits parabolic mass gain kinetics (i.e., passive oxidation with a protective scale) when  forms a homo is present,  phase,  it  Fig. 3.  Volatility diagram for SiC at 15001C after Heuer and Lou.30  Fig. 4.  Volatility diagram for ZrB2-SiC at 15001C.  \\x0c', '146  Journal of the American Ceramic Society—Fahrenholtz  Vol. 90, No. 1  (l)  the B2O3 passive oxidation with a non-protective  scale is stable, but  scale) when B2O3 evaporates. Likewise, ZrB2-SiC is expected to exhibit parabolic mass gain characteristics as long as the outer scale is stable.  linear mass gain kinetics  (i.e.,  (l)  Based on the relative vapor pressures of B2O3 (l) and SiO2 (l) and assuming ideal solution behavior, B2O3 (l) should evaporate preferentially. Thus, the stable scale on ZrB2-SiC at 15001C should consist, mainly, of SiO2 (l) and ZrO2 (cr). Thermodynamically, the combined ZrB2-SiC volatility diagram is consistent with experimental observations that SiC additions  improved the oxidation resistance of ZrB2-based ceramics above 12001C by forming a SiO2-rich protective scale.7,10,21,23-28 In the next section, the diagram is analyzed in more detail and exper imental observations of ZrB2-SiC oxidation are incorporated to understand the development of the SiC-depleted layer.  V.  Proposed Reaction Sequence for Depletion Layer  Formation  For ZrB2 (cr) and a-SiC (cr) individually, exposure to air at 15001C led to the formation of scales consistent with the pre dictions of their respective volatility diagrams. The response of  the combined system is more complex due to additional  inter actions beneath the outer scale. This paper represents the first  attempt to use a thermodynamic model to quantify the chemical  and physical  interactions that  lead to development of  the SiC depleted region. Four separate structures will be discussed:  (1)  unoxidized ZrB2-SiC; (2) the response during initial heating; (3) the structure as temperature approaches 15001C; and (4) the structure that develops after holding at 15001C. Chemical pro cesses that proceed during each step are specified, changes in the  physical  structure are discussed, and oxygen activities at each  interface are  calculated. This model  is  intended as a starting  point for more detailed experimental and modeling studies that  could focus on a particular layer,  interface, or chemical process  described in this analysis.  (1)  Unoxidized ZrB2-SiC  Initially, ZrB2-SiC consists of a ZrB2 (cr) matrix and dispersed a-SiC (cr) particles. A typical microstructure of ZrB2-SiC containing 30 vol% SiC particles is shown in the lower portion of  Fig. 1(a) and reproduced schematically in Fig. 5(a). No signifi cant solid solution is expected and ZrB2 and SiC are stable in contact with each other.  (2)  Initial Response During Heating  Thermodynamically, both ZrB2 and SiC should oxidize when exposed to air. Kinetically, the oxidation of ZrB2 has been reported to be more rapid than the oxidation of SiC below B12001C,19,24,26 which leads to the development of a structure  similar to that shown schematically in Fig. 5(b). The dominant  chemical process in this stage is the oxidation of ZrB2 by Reaction (1), which produces an oxide layer that consists of a con tinuous matrix of B2O3 (l) with entrained ZrO2 particles are also incorporated into the B2O3 (l) scale as the reaction layer penetrates into the ZrB2-SiC. Because the B2O3 (l) is continuous, it prevents direct exposure of ZrB2 to air and is, therefore, a protective scale below B12001C. At this stage, the  (cr). The SiC  SiC particles do not oxidize significantly and are not altered by  incorporation into the oxide.  (3)  Evolution as the Temperature Approaches 15001C  Approaching 15001C, the composition of the scale changes sig nificantly. Owing to its high vapor pressure, B2O3 (l) evaporates.18 Unlike the continuous film of B2O3 (l), ZrO2 (cr) does not form a protective scale.38 In this temperature range, the rate of  SiC oxidation increases, leading to the formation of SiO2 (l). The parabolic mass gain kinetics observed in this temperature regime  suggest a continuous evolution in the composition of the protective oxide layer from mainly B2O3 (l) below 12001C to mainly  Fig. 5.  Schematic diagram of a proposed reaction sequence for the for mation of the SiC-depleted layer during the oxidation of ZrB2-SiC at 15001C in air showing (a) unoxidized ZrB2-SiC, (b) the initial response during heating, (c) the structure when the specimen reaches 15001C, and (d) the steady-state response at 15001C.  (l) by 15001C rather  SiO2 than the complete loss of B2O3 (l), followed by the formation of a new SiO2 (l) layer.24 The ZrO2 (cr) that was formed initially remains concentrated near the  ZrB2-SiC surface, forming a ZrO2-SiO2 layer with a two-phase, interpenetrating, microstructure.23,26 The ZrO2-SiO2 layer has been reported to contain some B, although its distribution is  difficult to characterize due to the low sensitivity of energy-dispersive spectroscopy to light elements.26 It is likely that a small  fraction of B2O3 (l) remains dissolved in the SiO2 (l) in the ZrO2- SiO2 layer. The structure thought to exist at this stage is shown schematically in Fig. 5(c). The dominant chemical processes in stage are the oxidation of a-SiC (cr),  the evaporation of  this  B2O3 (l), and the formation of a coherent SiO2-rich scale.  (4)  Steady State at 15001C  (A)  15001C, ZrB2-SiC exhibits passive oxidation behavior with parabolic mass gain  Experimental Observations:  At  kinetics due to the formation of a protective scale (Fig. 5(d)).  The ZrO2-SiO2 layer remains between the outer SiO2 (l)-rich layer and the SiC-depleted ZrB2. Relatively thick (B100 mm or greater) ZrO2-SiO2 layers can be grown by thermal cycling26 or during arc heater testing,10,22 but in static oxidation the ZrO2- SiO2 layer remains thin compared with the SiC-depleted layer and the outer SiO2 (l) layer. Parabolic mass gain kinetics are observed because oxygen molecules or ions must diffuse through  both the SiO2 (l) layer and the underlying ZrO2-SiO2 layer to react with the ZrB2-SiC. If the ZrO2 (cr) or SiO2 (l) reacted with SiC (cr) or ZrB2 (cr), a non-protective scale and linear mass gain/loss kinetics would be expected (e.g., Reaction (4)). This  type of reaction seems unlikely as Reaction (4) and similar reenergies at 15001C and the  actions have positive Gibbs’  free  local structure does not  favor the rapid removal of any of  the  reaction products, which might allow the reaction to go forward despite its positive DG1. Opeka et al.1 suggested that  the tran sition to non-protective behavior  in ZrB2-SiC occurs over a  wide temperature range.  (B)  Oxygen Activity in the SiC-Depleted Region:  While  the pO2 relative rates of  in the SiC-depleted region may be  influenced by the  through the oxide scale (O2 source) and oxidation of SiC (O2 sink), its value is controlled by chemical reaction and its equilibrium value can,  therefore, be calcu transport  \\x0c', 'January 2007  Thermodynamic Analysis of ZrB2-SiC Oxidation  147  lated. Qualitatively,  the oxygen activity (pO2 should be significantly less than it is in the ambient atmosphere B2 \\x02 104 Pa in air) because of the chemical potential gra(pO2 dient associated with any diffusion profile. Based on experimen ) beneath the scale  tal observations, the oxygen activity below the scale is low enough to prevent oxidation of ZrB2 at 15001C as the ZrB2 in the SiC-depleted layer seems to remain unoxidized. The pO2 for ZrB2-ZrO2 equilibrium (Reaction (1)) at 15001C is 1.8 \\x02 10\\x0011 Pa, as discussed in the development of the ZrB2 volatility diain the SiC-depleted region  gram. This is an upper bound for pO2 and is well within the range of oxygen partial pressures  that  promote the active oxidation (i.e., formation of SiO (g) and CO (g)) of SiC at 15001C based on Fig. 3.  One way to estimate a lower bound for pO2 (g) stability. At low pO2 , CO (g) reduces to C (cr) at the so-called sooting limit (Reaction (5)). Solid C (cr) has not been reported in  is to consider CO  the SiC-depleted layer, so the pO2 most likely remains above the sooting limit. The nG1 for Reaction (5) at 15001C is 266 kJ. The  maximum pCO in the SiC-depleted layer can be read from Fig. 3 or calculated from the SiC-SiO2 equilibrium (Reaction (2)). Calculation gives a pCO of 1611 Pa at 15001C. Above this value, SiO2 (l) reacts with CO (g) to form SiC (reverse of Reaction (2)). Using pCO 5 1611 Pa, Reaction (5) was used to calculate a posof 5.0 \\x02 10\\x0015 Pa in the SiC-depleted sible lower bound of pO2 region.  CO ðgÞ ! C ðcrÞþ1  2O2  ðgÞ  (5)  Fig. 6.  Schematic showing the composition of  the layers  in oxidized  zirconium diboride (ZrB2)-SiC along with the relative interfacial motions, the chemical processes that are active, and the calculated oxygen  partial pressures.  (D)  Overall Process:  Based on the volatility diagrams  and the stability criteria defined by Reaction (5) and similar re actions, the structure depicted in Fig. 5(d) appears to be thermo dynamically stable. Even though the SiC-depleted region is created by active oxidation of a-SiC (cr) (linear mass loss kinet ics, no protective scale),  the overall process  is consistent with  passive oxidation behavior  and parabolic mass gain kinetics  based on the presence of a protective scale. Presumably, the rate limiting step of  the overall process  is  the diffusion of oxygen  through the SiO2-rich scale and the ZrO2-SiO2 layer. Based on this model, the thickness of the outer scale should continue to  (C)  Si Transport Across  the Depletion Layer:  For  the  increase with time due to the continued oxidation of SiC, which,  SiO2 (l) scale to grow by the mechanism proposed, silicon must be transported from ZrB2-SiC across the SiC-depleted layer to the oxide layer. Two possible driving forces for transport are a  in turn, increases the thickness of the SiC-depleted region in the  underlying ZrB2-SiC. Somehow, the oxide scale remains coherent and attached to  temperature gradient and a chemical potential gradient. As iso the surface despite the formation of voids in the SiC-depleted  thermal oxidation is assumed in this study, only the latter will be  region.  In contrast,  formation of voids at  the scale-substrate  considered as a driving force for SiO (g)  transport across  the  SiC-depleted layer. One possibility is that pO2 an isothermal SiC-depleted region and that the concentration of  is constant across  some other species drives SiO (g) transport. Arbitrarily, if the were set at the upper limit, 1.8 \\x02 10 \\x0011 Pa, then SiC value of pO2 in the ZrB2-SiC would oxidize and a layer of SiO2 (l) would form on its surface. The vapor pressure of SiO (g) above SiO2 (l) is high at this pO2 , approximately 100 Pa from Fig. 3 (344 Pa by calculation). For an isothermal system with a uniform pO2 pSiO would then be identical on both sides of the SiC-depleted layer as both would consist of SiO2 (l), which would not drive SiO (g) transport in either direction. Therefore, a pO2 gradient is likely to exist across the SiC-depleted layer.  , the  From the above discussion, the pO2 at the interface between the SiC-depleted layer and the ZrB2-SiC must be o1.8 \\x02 10\\x0011 Pa to drive SiO (g) transport. If the pO2 at this interface were above the pO2 for SiC-SiO2 equilibrium (Reaction (2)), which is 8.7 \\x02 10\\x0013 Pa, then a layer of SiO2 (l) would be expected to form on the surface of the ZrB2-SiC layer. Even this slight pO2 gradient would drive SiO (g) transport across the SiC-depleted  de pO2  region because of  the increase in pSiO that occurs as pO2 creases for SiO2 (l), as shown in Fig. 3. If pO2 were lower than the value for SiC-SiO2 equilibrium, then the surface of the ZrB2- SiC layer would be free of SiO2 (l). For SiC, the pSiO decreases as decreases, which allows a minimum threshold pO2 to be defined based on the fact that pSiO must be higher at the interface between the SiC-depleted region and the ZrB2-SiC than at the other interface. Using the upper pO2 limit in the SiC-depleted layer, Reaction (4) can be used to calculate pSiO 5 344 Pa. Given that pSiO must be higher at the surface of the ZrB2-SiC than at the other interface, the minimum pO2 value at the ZrB2-SiC surface can be calculated by considering the SiC-SiO equilib rium (Reaction (3)). Setting pSiO 5 344 Pa, the minimum pO2 at the ZrB2-SiC surface can be calculated as 4.1 \\x02 10 \\x0014 Pa. The calculated pO2 values and probable chemical processes active during oxidation at 15001C in air are summarized schematically  in Fig. 6.  interface in pure SiC disrupts the scale, causing the loss of passive oxidation protection.35,39 As a consequence of the pressure  of CO (g) and SiO (g) beneath the SiO2 (l) scale, some threshold temperature is likely to exist at which ZrB2-SiC will exhibit nonprotective oxidation behavior. As the temperature increases, the  pressure in the SiC-depleted region will  increase until  it is high  enough to rupture the scale or cause failure at one of the inter faces in the layered structure. This may occur when the total pressure in the SiC-depleted region reaches B1.013 \\x02 105 Pa (1 atm),1 which is estimated to occur at B17751C. The scale has  the potential to be self-healing in the sense that the local oxida tion of SiO (g) or flow of the viscous outer scale could quickly  re-establish protective behavior if  the coating were comprom ised. Experimentally, arc heater testing of ZrB2-SiC and HfB2- SiC has shown that these materials exhibit protective oxidation  behavior at temperatures above the stability limit of SiO2 due to the formation of a protective ZrO2 layer.  VI.  Summary and Conclusions  The oxidation of ZrB2-SiC in air at 15001C was analyzed using volatility diagrams for ZrB2 and SiC. The diagrams justified, the development of a SiC-depleted region  thermodynamically,  beneath the outer oxide scale. Owing to the high vapor pressure  of SiO (g) under  reducing conditions, active oxidation of SiC  under the oxide scale produced a SiC-depleted region. The pO2 in the SiC-depleted layer was calculated to be between 4.1 \\x02 10\\x0014 and 1.8 \\x02 10\\x0011 Pa. The outer SiO2 scale grows by transport of SiO (g) from the unoxidized ZrB2-SiC to the scale due to a pO2 gradient. The proposed process is consistent with passive oxi dation and parabolic mass gain kinetics as a protective scale  forms and continues to increase in thickness during oxidation.  The formation of a SiC-depleted region appears inevitable when ZrB2-SiC is heated in air at 15001C due to active oxidation of SiC. Because SiC is the source of SiO for growth of the  outer scale, the thickness of the SiC-depleted layer, most likely,  \\x0c', 'continues  to increase as  the  thickness of  the  scale  increases.  Based on the formation of the SiC-depleted region, ZrB2-SiC may not be usable for extended periods above 17751C in air due  to gas pressure beneath the scale. However, ZrB2-SiC has demonstrated the remarkable ability to withstand oxidizing condi tions over a wide temperature range.  Acknowledgments  The author would like to thank Dr. Jeff Smith of UMR and Dr. Mark Opeka  of  the Naval Surface Warfare Center-Carderock Division for many helpful dis cussions. 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Opeka, ‘‘Properties of Ceramics in the  ZrB2/ZrC/SiC System Prepared by Reactive Processing,’’ Ceram. Eng. Sci. Proc., 19 [3] 105-12 (1998). 39D. Das, J. Farjas, and P. Roura,  ‘‘Passive-Oxidation Kinetics in SiC Micro particles,’’ J. Am. Ceram. Soc., 87 [7] 1301-5 (2004).  &  148  Journal of the American Ceramic Society—Fahrenholtz  Vol. 90, No. 1  \\x0c']"
},{
  "_id": 272,
  "PDF": "Thermodynamic Calculation of HfB2 Volatility Diagram.pdf",
  "Text": "['Section I: Basic and Applied Research  JPEDAV (2011) 32:422-427 DOI: 10.1007/s11669-011-9930-x 1547-7037 ÓASM International  Th e rmod yn am ic C a lcu la t ion o f H fB 2 Vo la t i l i t y D iag ram  H u i L i , L i tong Z ha n g , Q i ng f en g Ze n g , a nd La i fe i C h en g  (Submitted January 18, 2011;  in revised form May 20, 2011)  The thermodynamics of the oxidation of HfB2 at temperatures of 1000, 1500, 2000, and 2500 K have been studied using volatility diagrams. Both the equilibrium oxygen partial pressure (PO2 ) for the HfB2(s) to HfO2(s) plus B2O3(l) and the partial pressures of B-O vapor species formed due to B2O3(l) volatilization increase with increasing temperature. Vapor pressures of the predominant gaseous species also increase with PO2 . At 1000 K, the predominant vapor transition sequence is predicted be BO(g) fi B2O2(g) fi B2O3(g) fi BO2(g) with increasing PO2 , and the 1.2731026 Pa predominant gas is BO2(g) with a pressure of under the condition of PO2 = 20 kPa. At higher temperatures of 1500, 2000, and 2500 K, the system undergoes vapor transitions in the same sequence of B(g) fi BO(g) fi B2O2(g) fi B2O3(g) fi BO2(g). Under the same condition of PO2 = 20 kPa, the predominant vapor species is B2O3(g) with pressures of 2.38, 4.493103, and 3.553105 Pa, respectively. Volatilization of B2O3(l) may produce porous HfO2 scale, which is consistent with the experimental observations of HfB2 oxidation in air. The present volatility diagram of HfB2 shows that HfB2 exhibits oxidation behavior similar to ZrB2, and factors other than volatility of gaseous species affect the oxidation rate.  Keywords  FactSage, HfB2, oxidation thermodynamics, volatility diagram  1 .  In t rodu c t ion  Refractory transition metal borides, known as ultra high temperature ceramics (UHTCs), are candidates for applications in extreme thermal and chemical environments.[1-8] Among this material family, hafnium diboride (HfB2) is of particular interest because it has an unusual combination of physical properties including a sufficiently high melting point and hardness, retained strength at high temperature, modest coefficient of thermal expansion, as well as the ability to form a refractory oxide scale. These features give HfB2 the capability to withstand temperatures in the range of 2173-2773 K in oxidizing environments.[1] A review of studies on the oxidation of the UHTCs has been provided in the comprehensive work of Opeka et al.[2] The oxidation of HfB2 in air produces an oxide scale composed of HfO2 and B2O3. The oxidation layer of HfO2 is a porous scale, and the distribution of HfO2 and B2O3 in the oxide scale varies with (<\\x181273 K), temperatures.[8-13] At low temperatures a liquid B2O3 film is formed on the top of HfO2 plus B2O3 scale, and the pores in the HfO2 are covered or filled with glassy or liquid B2O3. At intermediate temperatures (\\x181273 to \\x182073 K), B2O3(l) would still partially fill the pores in porous HfO2 even though the external B2O3(l) layer evaporates. The remaining B2O3(l) layer acts as a barrier to oxygen diffusion, resulting in passive oxidation with  Hui Li, Litong Zhang, Qingfeng Zeng, and Laifei Cheng, National  Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, Xi’an 710072, Shaanxi, People’s Republic of China. Contact e-mail: huili128@gmail.com.  ers[15,16]  para-linear oxidation kinetics.[10,11] At high temperatures (>\\x182073 K), B2O3(l) would be absent owing to its volatilization, yielding a non-protective porous HfO2(s) layer. Oxidation is a dynamics process, sensitive to perturbations in temperature and oxidant concentration. A volatility diagram[2,14-19] plots the vapor pressure of the predominant gaseous species in equilibrium with the various condensed phases as a function of oxygen partial pressure and temperature, so it is appropriate for understanding the oxidation behaviors of UHTCs when more than one gaseous species is involved. Thermodynamic calculations have been conducted on the volatility diagrams by Lou and coworkto study the gas-solid interactions for material systems such as Mg-O, Al-O, Si-O, Si-N, Si-N-O and Si-CO. In addition, volatility diagrams have been used to reveal for B,[2] significant insight into the oxidation behaviors and ZrSiO4 in extremely high-temperature environments. However, up to now, the volatility diagram for HfB2 has not been reported. The oxidation behaviors of ZrB2 and HfB2, perhaps analogous, will be different to some extent. It was shown that the oxidation resistance of HfB2 was superior to that of ZrB2 from 1473 to Therefore, the volatility diagram may provide insight into the oxidation mechanism of pure HfB2 and serve as the basis for the investigation of the oxidation behavior of HfB2-based including HfB2-SiC.[2,11,20,21] The ceramics constructions of volatility diagrams for HfB2 varying with temperature and equilibrium oxygen pressure are described in this paper.  Zr,[2,19] ZrB2,[14] SiC,[17] ZrB2-SiC,[18]  [19]  2473 K.[2,8-11,20,21]  2 . Compu t a t iona l M e th od  The dominant species at the solid-gas interface are determined according to the principle of chemical equilib 422  Journal of Phase Equilibria and Diffusion Vol. 32 No. 5 2011  \\x0c', 'rium by minimizing the Gibbs free energy of the system.[22-24] The present thermodynamic calculations for the construction of HfB2 volatility diagram are done by means of the FactSageä software package.[24] The thermodynamic data are used from the database FACT53 (version 5.3). All present thermodynamic calculations are made at standard state, with the constraints of closed system and mass conservation. It is assumed that pure HfB2 oxidizes in the dry air condition (no water is present) with an ambient pressure of 1.019105 Pa. At given different temperatures (1000, 1500, 2000, and 2500 K), HfB2 will oxidize to a series of solid and gaseous products according to the related chemical reactions. Then with the assumption of unit activity for the condensed phases HfB2(s), HfO2(s) and B2O3(l), the corresponding equilibrium oxygen partial pressure and the amounts of gaseous species will be determined. Reactions among gaseous species are not considered. Finally the volatility diagram will be constructed.  3 . Re su l t s and D i scu s s ion  3 .1 Co n s t ruc t io n o f H fB 2 Vo l a t i l i t y D i ag ram  The reactants and products involved in the reactions for the construction of HfB2 volatility diagram are listed in Table 1. When HfB2 is oxidized above the melting temper(\\x18723 K[20]), ature of B2O3 an oxide scale consisting of HfO2 and B2O3 is formed according to Reaction 1.  [9,12,13]  HfB2 ðsÞ þ 5 2  O2 ðgÞ ! HfO2 ðsÞ þ B2O3 ðlÞ  ðEq 1Þ  The change in Gibbs free energy (DrG\\x0e ) for Reaction 1 is firstly calculated, and then converted to reaction equilibrium constant (Keq ) value using Eq 2:  DrG\\x0e ¼ \\x00RT ln Keq  ðEq 2Þ  where R is temperature. In Reaction 1, unit activity is assumed for the condensed phases HfB2(s), HfO2(s) and B2O3(l), and thus the corresponding equilibrium oxygen partial pressure (PO2 ) for oxidation of HfB2 is determined based on the following relation:  the  ideal gas  constant  and T  is  the  absolute  Keq ¼ ðaB2 O3 ðlÞ ÞðaHfO2 ðsÞ Þ ðaHfO2 ðsÞ ÞðaO2 ðgÞ Þ5=2 ¼  1  ðPO2 Þ5=2  ðEq 3Þ  where a is the activity of species involved in the reaction. At 1500 K with ambient pressure of 1.019105 Pa, Keq is 1.0991050 for Reaction 1, and then the value of equilibrium 9.80910\\x0016 Pa (log PO2 = \\x0015.01) of for the co-existence of HfB2(s) and HfO2(s) plus B2O3(l) is obtained using Eq 3. The vertical line at the equilibrium separates the diagram into two regions as shown in Fig. 1. Below the equilibrium PO2 (left side of diagram), the vapor species are in equilibrium with HfB2(s). For example, B2O3(g) is formed by Reaction 4. Then, the corresponding reaction equilibrium constant and equilibrium partial pressure of B2O3(g) can be determined using Eq 5.  PO2  PO2  HfB2 ðsÞ þ 5 2  O2 ðgÞ ! HfO2 ðsÞ þ B2O3 ðgÞ  ðEq 4Þ  Keq ¼ ðPB2 O3 Þ=ðPO2 Þ5=2 ;  logðPB2 O3 Þ ¼ logðKeq Þ þ 5 2  logðPO2 Þ  ðEq 5Þ  As shown in Eq 5, the logarithmic amount of B2O3(g) should linearly increase with logarithmic PO2 since oxygen is a reactant. In this region, other gaseous B-O and Hf-O species formed by the reactions of HfB2 using the same methodology are listed in Table 2. Above the equilibrium  Table 1  Hf, B, and O species for calculation of  the HfB2 volatility diagram  Elements  Hf  B  O  Hf-B  Hf-O  B-O  Gas species  Hf(g)  B(g)  O2(g)  HfO(g)  BO(g) BO2(g) B2O(g)  B2(g)  HfO2(g)  B2O2(g) B2O3(g)  Condensed species  HfB2(s)  HfO2(l)  HfO2(s1, s2)  B2O3(l)  −30 −35  −30  −25  −20  −15  −10  −5  0  5  10  −25  −20  −15  −10  −5  0  5  10  log P O2  /Pa  l  g o  P  (  x  )  /  P  a  B(g)  B2(g)  BO(g)  B2O(g)  BO2(g)  B2O2(g)  B2O3(g)  Hf(g)  HfO(g)  HfO2(g)  HfB2(s)  HfO2(s)+B2O3(l)  Fig. 1  Partial pressures of B-O and Hf vapor species as a function of oxygen partial pressure at 1500 K  Basic and Applied Research: Section I  Journal of Phase Equilibria and Diffusion Vol. 32 No. 5 2011  423        \\x0c', 'Section I: Basic and Applied Research  Table 2 Volatilization reactions and their corresponding temperature-dependent vapor partial pressures involving HfB2(s) as the primary condensed phase  Reactions producing volatile B or Hf species  HfB2 ðsÞ þ 5 2 O2 ðgÞ ! HfO2 ðsÞ þ B2O3 ðgÞ HfB2 ðsÞ þ 2O2 ðgÞ ! HfO2 ðsÞ þ B2O2 ðgÞ HfB2 ðsÞ þ 3 2 O2 ðgÞ ! HfO2 ðsÞ þ B2OðgÞ HfB2 ðsÞ þ O2 ðgÞ ! HfO2 ðsÞ þ B2 ðgÞ HfB2 ðsÞ þ 3O2 ðgÞ ! HfO2 ðsÞ þ 2BO2 ðgÞ HfB2 ðsÞ þ 2O2 ðgÞ ! HfO2 ðsÞ þ 2BOðgÞ HfB2 ðsÞ þ O2 ðgÞ ! HfO2 ðsÞ þ 2BðgÞ HfB2 ðsÞ þ 5 2 O2 ðgÞ ! HfO2 ðgÞ þ B2O3 ðlÞ HfB2 ðsÞ þ 2O2 ðgÞ ! HfOðgÞ þ B2O3 ðlÞ HfB2 ðsÞ þ 3 2 O2 ðgÞ ! Hf ðgÞ þ B2O3 ðlÞ  Vapor pressure  log PB2 O3 ¼ logðKeq ðT ÞÞ þ 5 2 logðPO2 Þ log PB2 O2 ¼ logðKeq ðT ÞÞ þ 2 logðPO2 Þ log PB2 O ¼ logðKeq ðT ÞÞ þ 3 2 logðPO2 Þ log PB2 ¼ logðKeq ðT ÞÞ þ logðPO2 Þ log PBO2 ¼ 1 2 ½logðKeq ðT ÞÞ þ 3 logðPO2 Þ\\x8a log PBO ¼ 1 2 ½logðKeq ðT ÞÞ þ 2 logðPO2 Þ\\x8a log PB ¼ 1 2 ½logðKeq ðT ÞÞ þ logðPO2 Þ\\x8a log PHfO2 ¼ logðKeq ðT ÞÞ þ 5 2 logðPO2 Þ\\x8a log PHfO ¼ logðKeq ðT ÞÞ þ 2 logðPO2 Þ\\x8a log PHf ¼ logðKeq ðT ÞÞ þ 3 2 logðPO2 Þ\\x8a  Table 3 Volatilization reactions and their corresponding temperature-dependent vapor partial pressures involving HfO2(s) and B2O3(l) as the primary condensed phase  Reactions producing volatile B or Hf species  B2O3 ðlÞ ! B2O3 ðgÞ B2O3 ðlÞ ! B2O2 ðgÞ þ 1 2 O2 ðgÞ B2O3 ðlÞ ! B2OðgÞ þ O2 ðgÞ B2O3 ðlÞ ! B2 ðgÞ þ 3 2 O2 ðgÞ B2O3 ðlÞ þ 1 2 O2 ðgÞ ! 2 BO2 ðgÞ B2O3 ðlÞ ! 2 BOðgÞ þ 1 2 O2 ðgÞ B2O3 ðlÞ ! 2BðgÞ þ 3 2 O2 ðgÞ HfO2 ðsÞ ! HfO2 ðgÞ HfO2 ðsÞ ! HfOðgÞ þ 1 2 O2 ðgÞ HfO2 ðsÞ ! Hf ðgÞ þ O2 ðgÞ  PO2  (right side of Fig. 1), the gaseous species are in equilibrium with HfO2(s) and B2O3(l). In this regime, B2O3(l) directly vaporizes to B2O3(g) according to Reaction 6, and the corresponding partial pressure is determined by Eq 7: B2O3 ðlÞ ! B2O3 ðgÞ  ðEq 6Þ  PB2 O2  Keq ¼ ðPB2 O3 Þ;  logðPB2 O3 Þ ¼ logðKeq Þ  ðEq 7Þ  Because oxygen is neither consumed nor produced by Reaction 6, PB2O3 does not vary with PO2 in this regime, resulting in a horizontal line. Similarly, other lines in this regime are plotted according to the reactions listed in Table 3. At the equilibrium PO2 of Reaction 1, the partial pressure of each gaseous specie above HfB2(s) must be the same as it is above HfO2(s)-B2O3(l) since the gases are in equilibrium with all of the condensed phases at this PO2 . Vapor pressures of all the gaseous species were calculated at 1500 K in this manner as shown in Fig. 1. By omitting the species with lower pressures and showing only the predominant species identified in Fig. 1, the volatility diagram for HfB2 was constructed at 1500 K (Fig. 2). Volatility diagrams for HfB2 were also completed at 1000, 2000, and 2500 K by employing the same methodology, as shown in Fig. 3.  Vapor pressure  logðPB2 O3 Þ ¼ logðKeq ðT ÞÞ logðPB2 O2 Þ ¼ logðKeq ðT ÞÞ \\x00 1 2 logðPO2 Þ logðPB2 O Þ ¼ logðKeq ðT ÞÞ \\x00 1 logðPO2 Þ logðPB2 Þ ¼ logðKeq ðT ÞÞ \\x00 3 2 logðPO2 Þ logðPBO2 Þ ¼ 1 2 ½logðKeq ðT ÞÞ þ 1 2 logðPO2 Þ\\x8a logðPBO Þ ¼ 1 2 ½logðKeq ðT ÞÞ \\x00 1 2 logðPO2 Þ\\x8a logðPB Þ ¼ 1 2 ½logðKeq ðT ÞÞ \\x00 3 2 logðPO2 Þ\\x8a logðPHfO2 Þ ¼ logðKeq ðT Þ) logðPHfO Þ ¼ logðKeq ðT ÞÞ \\x00 1 2 logðPO2 Þ logðPHf Þ ¼ logðKeq ðT ÞÞ \\x00 logðPO2 Þ  a  P  /  )  x  (  P  g o  l  10  5  0  −5  −10  −15  −20 −35  HfB2(s)  HfO2(s)+B2O3(l)  B(g)  BO(g)  B2O2(g)  B2O3(g)  BO2(g)  0.45  0.40  a  P  /  )  x  (  P  g o  l  0.35 −15.09  −14.99  log PO  2   /Pa  −14.89  −30  −25  −20  −15  −10  −5  0  5  10  log PO2 /Pa  Fig. 2  Volatility diagrams for  the oxidation of HfB2 at 1500 K  3 .2 D i sc u s s ion o f H fB 2 Vo l a t i l i t y D i ag r am  Volatility diagrams are isothermal plots showing the partial pressures of gaseous species in equilibrium with the possible condensed phases in a system. Figure 1 shows that the stable solid phase changes from HfB2 to HfO2 at around logPO2 = \\x0015.01 Pa.  424  Journal of Phase Equilibria and Diffusion Vol. 32 No. 5 2011          \\x0c', 'Basic and Applied Research: Section I  Table 4 Pressures of gaseous products for HfB2 oxidation in air (PO2 = 20 kPa) as a function of temperature at 1000, 1500, 2000, and 2500 K  Vapor pressures, Pa  Species  1000 K  1500 K  2000 K  2500 K  BO2(g)  B2O3(g)  B2O2(g)  BO(g)  B2O(g)  B2(g)  B(g)  HfO2(g)  HfO(g)  Hf(g)  1.27910\\x006 3.71910\\x007 2.02910\\x0023 2.96910\\x0017 7.32910\\x0048 1.36910\\x0081 2.95910\\x0043 1.00910\\x0043 4.18910\\x0040 1.49910\\x0068  7.37910\\x001  2.38 5.53910\\x0010 1.25910\\x006 6.30910\\x0025 8.92910\\x0046 8.48910\\x0023 3.40910\\x0024 8.24910\\x0021 1.06910\\x0038  4.889102  4.499103 2.14910\\x003 2.09910\\x001 1.25910\\x0013 5.03910\\x0028 1.21910\\x0012 1.47910\\x0014 2.67910\\x0011 7.29910\\x0024  2.229104  3.559105  1.639101  2.549102 6.05910\\x007 1.83910\\x0017 1.36910\\x006 6.09910\\x009 9.24910\\x006 4.29910\\x0015  10  5  0  −5  −10  −15  a  P  /  )  x  (  P  g o  l  HfB2(s)  HfO2(s)+B2O3(l)  2500 K  2000 K  1500 K  1000 K  B(g)  BO(g)  B2O2(g)  B2O3(g)  BO2(g)  −20 −35  −30  −25  −20  −15  −10  −5  0  5  10  log P O2 /Pa  Fig. 3  Volatility diagrams of HfB2 oxidation as temperature at 1000, 1500, 2000, and 2500 K  a  function of  PO2  As shown in Fig. 1, the highest vapor pressures for several gaseous species exist at the HfB2-oxide interface, but BO2(g) is an exception. Its pressure increases as PO2 increases because the formation of most of the B-containing or Hf-containing gases produce oxygen due to the decomposition reaction of B2O3(l) or HfO2(s) whereas oxygen is consumed for the formation of BO2(g). As a result, the logarithm of the partial pressure of BO2(g) increases in direct proportion to the logarithm of the partial pressure of O2. HfB2 volatility diagrams (Fig. 2, 3) show the equilibrium oxygen partial pressures for the HfB2(s) to HfO2(s) plus B2O3(l) as well as the variations of predominant gaseous products with changing and temperature. Figure 2 exhibits that five predominant vapor phases are present for the oxidation of HfB2 at 1500 K. The system underin sequence of B(g) fi BO(g) fi goes vapor transitions B2O2(g) fi B2O3(g) fi BO2(g) with increasing PO2 , and the values of PO2 at the four vapor transition positions are 9.32910\\x0029, 3.28910\\x0027, 1.10910\\x0015, and 2.219106 Pa, respectively. Volatility diagrams of HfB2 at 2000 and 2500 K are similar with that at 1500 K, and they undergo the same vapor transitions sequence (see Fig. 3). An apparent difference at 1000 K is that the plot does not show the vapor transition from B(g) to BO(g) since this occurs below the limit of the plot. As a result, the vapor transitions at 1000 K appear to be BO(g) fi B2O2(g) fi B2O3(g) fi BO2(g) with increasing PO2 . In addition, as the trend shown in Fig. 3, both the partial pressures of B-O vapor species due to B2O3(l) evaporation and the equilibrium oxygen partial pressures for the HfB2(s) to HfO2(s) plus B2O3(l) increase with temperature. For example, at temperatures of 1000, 1500, 2000, and 2500 K, the condensed phase transition from the HfB2(s) to HfO2(s) plus B2O3(l) occurs at around log PO2 = \\x0028.96, \\x0015.01, \\x008.11, and \\x004.04 Pa, respectively. Volatility diagrams can be used to interpret experimental observations of HfB2 oxidation in air. At a typical oxidizing environment of PO2 = 20 kPa, pressures of all vapor species for the oxidation of HfB2 are computed as a function of  tions[2,9,12,13]  temperature, and the results are summarized in Table 4. The volatilization rate for B2O3(l) in air is expected to be low at 1000 K based on the partial pressures of the various gases above HfO2(s) plus B2O3(l) layer, the maximum of which is only 1.27910\\x006 Pa. The low partial pressures of various gases are consistent with the experimental observathat a liquid B2O3 film is formed on top of the HfO2 and B2O3 scale at low temperatures. At 1500 K, the predominant vapor species is B2O3(g) with a pressure of 2.38 Pa. This value is about six orders of magnitude larger than that at 1000 K, indicating that the B2O3(l) vaporizes at 1500 K, which maybe the reason for the disappearance of the external B2O3 layer and some of the B2O3(l) in the pores. At higher temperatures of 2000 K and 2500 K, the predominant vapor pressures of B2O3(g) are 4.499103 and 3.559105 Pa, respectively. It is also shown in Table 4 that, at an oxygen partial pressure of 20 kPa, the vapor pressures of the B-O gaseous species at 2500 K are at least ten orders of magnitude higher than those at 1000 K. The oxidation of HfB2 studied at temperatures of 1480-2000 K revealed that the HfO2(s) should be the only condensed oxide on the vaporize.[10] surface of HfB2, as B2O3(l) should The calculated high partial pressures of vapor species at 1500, 2000, and 2500 K might result in rapid volatilization of B2O3(l), which is in accordance with the previous experimental study. Moreover, considering the units conversion, Table 4 shows that the pressures of B-containing gaseous products of HfB2 from 1000 to 2500 K are on the same order of magnitude with those of ZrB2 reported by Fahrenholtz.[14] The small deviations may stem from differences in the data or rounding during calculations. By calculations, the vapor pressures of HfO2(g), HfO(g) and Hf(g) are at least one order of magnitude lower than those of ZrO2(g), ZrO(g) and Zr(g) species. Based on the calculations, the vapor pressures of the Hf-O species should be lower than Zr-O species. From the thermodynamic, the B-O species are predominant in both systems, which would indicate that the oxidation behaviors of HfB2 and ZrB2 should be similar based on the present volatility diagram of HfB2. However, the relevant indicates that HfB2 is more oxidation than ZrB2, which  literature[2,8-11,20,21]  Journal of Phase Equilibria and Diffusion Vol. 32 No. 5 2011  425  resistant      \\x0c', 'Section I: Basic and Applied Research  indicates that factors other than volatility of gaseous species affect the oxidation rate. The calculated B2O3(g) interfacial vapor pressure at HfB2-HfO2 interface exceeds 1.019105 Pa at around 2331 K. Volatilization of B2O3(l) leaves behind a porous and non-protective HfO2 scale. Additives that reduce the activity of B2O3(l), such as SiC or other SiO2-formers, may improve the oxidation resistance of HfB2 at high temperatures by reducing the vapor pressure of B-based species. More detailed analysis is required to understand the complex interactions in these systems.  4 . Con c lu s ion s  Volatility diagrams for the oxidation of HfB2 have been constructed at temperatures of 1000, 1500, 2000, and 2500 K. These volatility diagrams reveal equilibrium partial oxygen pressures for the transition from HfB2(s) to HfO2(s) and B2O3(l) as a function of temperature, as well as vapor pressures of the predominant vapor species as a function of temperature and PO2 . Volatility diagram calculations are helpful to understand the oxidation behaviors of HfB2. The main conclusions are summarized as follows:  (1)  (2)  (3)  The equilibrium PO2 for the HfB2(s) to HfO2(s) plus B2O3(l) increase with temperature. For example, at temperatures of 1000, 1500, 2000, and 2500 K, the equilibrium oxygen partial pressures log PO2 are \\x0028.96, \\x0015.01, \\x008.11, and \\x004.04 Pa, respectively. Vapor pressures of the predominant gaseous species increase with both temperature and PO2 . At 1000 K, the predominant vapor transition sequence is predicted be BO(g) fi B2O2(g) fi B2O3(g) fi BO2(g) with increasing PO2 , and the predominant gas is BO2(g) with pressure of 1.27910\\x006 Pa under the condition of PO2 = 20 kPa. Increasing the temperatures from 1500 to 2000 K and 2500 K, the system undergoes vapor transitions in the same sequence of B(g) fi BO(g) fi B2O2(g) fi B2O3(g) fi BO2(g), and the predominant vapor species is B2O3(g) with the pressures of 2.38, 4.499103, and 3.559105 Pa, respectively. From the thermodynamic viewpoint at standard state, although the vapor pressures of the Hf-O species are lower than the corresponding Zr-O species under similar conditions, volatility diagram calculations show that HfB2 exhibits the similar oxidation behavior to ZrB2. The literature report that HfB2 is more oxidation resistant than ZrB2 indicates that factors other than volatility of gaseous species affect the oxidation rate.  Ac know l edgm en t s  This work is Key Laboratory  supported by Research Fund of the State of Solidification Processing of NWPU,  11.  426  Journal of Phase Equilibria and Diffusion Vol. 32 No. 5 2011  China (Grant No. 65-TP-2011), the Natural Science Foundation of China (Grant No. 50802076), and the 111 Project (B08040). The authors also acknowledge the Northwestern Polytechnical University High Performance Computing Center for the allocation of computing time on their machines.  Re f e re nc es  1. S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, and J.A. Salem, Evaluation of Ultra-High Temperature Ceramics for Aeropropulsion Use, J. Eur. Ceram. Soc., 2002, 22(14-15), p 2757-2767 2. M.M. Opeka, I.G. Talmy, and J.A. Zaykosk, Oxidation-Based Materials Selection for 2000 °C + Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience, J. Mater. Sci., 2004, 39(19), p 5887-5904 3. R. Savino, M.D.S. Fumo, D. Paterna, and M. 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Fahrenholtz, The ZrB2 Volatility Diagram, Ceram. Soc., 2005, 88(12), p 3509-3512 15. V.L.K. Lou, T.E. Mitchell, and A.H. Heuer, Review—Graphical Displays of the Thermodynamics Of high-Temperature Gas-Solid Reactions and Their Application to Oxidation of Metals and Evaporation of Oxides, J. Am. Ceram. Soc., 1985, 68(2), p 49-58  III. Soc.,  10.  J.  Am.  \\x0c', 'Basic and Applied Research: Section I  16. A.H. Heuer and V.L.K. Lou, Volatility Diagrams for Silica, Silicon Nitride, and Silicon Carbide and Their Application to High-Temperature Decomposition and Oxidation, J. Am. Ceram. Soc., 1990, 73(10), p 2789-2803 17. T. Goto, High-Temperature Oxidation Behavior of Chemical Vapor Deposited Silicon Carbide, J. Ceram. Soc. Jpn., 2002, 110(10), p 884-889 18. W.G. Fahrenholtz, Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region, J. Am. Ceram. Soc., 2007, 90(1), p 143-148 19. P. Barreiro, P. Rey, A. Souto, and F. Guitia´ n, Porous Stabilized Zirconia Coatings on Zircon Using Volatility Diagrams, J. Eur. Ceram. Soc., 2009, 29(4), p 653-659 20. C.M. Carney, Oxidation Resistance of Hafnium DiborideSilicon Carbide from 1400 to 2000 °C, J. Mater. Sci., 2009, 44(20), p 5673-5681  22.  21. C.M. Carney, T.A. Parthasarathy, and M.K. Cinibulk, Oxidation Resistance of Hafnium Diboride Ceramics with Additions of Silicon Carbide and Tungsten Boride or Tungsten Carbide, J. Am. Ceram. Soc., 2011. doi:10.1111/j.1551-2916.2011.04462.x J.J. Wang, L.T. Zhang, Q.F. Zeng, L.F. Cheng, and Y.D. Xu, Modified Wagner Model for the Active-to-Passive Transition in the Oxidation of Si3N4, J. Phys. D, 2008, 41(11), p 115412 J.J. Wang, L.T. Zhang, Q.F. Zeng, G.L. Vignoles, and A. Guette, Theoretical Investigation for the Active-to-Passive Transition in the Oxidation of Silicon Carbide, J. Am. Ceram. Soc., 2008, 91(5), p 1665-1673 24. C. Bale, P. Chartrand, S.A. Degterov, G. Eriksson, K. Hack, R. Ben Mahfoud, J. 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},{
  "_id": 273,
  "PDF": "Toward Oxidation-Resistant ZrB2-SiC Ultra High Temperature Ceramics.pdf",
  "Text": "['Toward Oxidation-Resistant ZrB2-SiC Ultra High Temperature Ceramics  EMILY EAKINS, DONI DANIEL JAYASEELAN, and WILLIAM EDWARD LEE  Ultra  high temperature  ceramics  (UHTCs),  including ZrB2-SiC, are designed for extreme in which temperatures exceed 2273 K (2000 °C). A key material their resistance to oxidation. Recent research  environment  applications  property of UHTCs  in many applications  is  into UHTCs  is described,  revealing  a  variety of different methods  for  improving  the oxi dation performance, which include  control of  starting powders,  composition and size dis tribution, mixing, and densification techniques. The use of additives has also been researched  widely,  for  example,  to increase  the  viscosity of  any  liquid phase  formed or provide pro tective  refractory phases  at high temperatures. SiC additions are effective in forming pro~1873 K (1600 °C). formation of a dense ZrO2 produce self-generating  tective  silica  but  only  in  static  environments  and  to  For  higher  temperature  applications,  additions  of La  lead  to  the  scale  probably  via  liquid  phase  sintering.  Such  ceramic  systems, which  refractory oxidation barriers or dense ZrO2 oxidation-resistant UHTCs.  scales,  show the greatest promise  in providing  DOI: 10.1007/s11661-010-0540-8 Ó The Minerals, Metals & Materials Society and ASM International 2010  I.  INTRODUCTION  ULTRA high designed to operate  temperature  ceramics  (UHTCs)  are  in extreme  environments  such as  those  experienced  on  leading  edges  of  hypersonic  vehicles or  in propulsion components of missiles.  In  recent years,  there has been a resurgence of  interest  in  UHTCs, particularly methods of  improving their high temperature capabilities.  For UHTCs  to maintain  their  structural  integrity  during service, exceptional oxidation resistance is par amount. Diborides  of  the  transition metals  such  as  hafnium and zirconium have proven to be some of  the  best candidate UHTC materials to date. For this review,  only zirconium diboride-based ceramics will be consid ered  unless  research  into  hafnium diboride  provides  useful  insight.  Several  reviews  have  summarized  the  history of UHTC research, material testing methods,[1,2]  properties,  and  and  these  reviews  consider  the  candidate materials and selection procedures necessary  to produce a UHTC that fulfils the stringent requireenvironment applications.[3-5]  ments  for  these  extreme  The purpose of this review is to summarize the advances  in research conducted in recent years  that  identify the  possible methods of  improving the oxidation resistance  of existing UHTC materials.  II.  OXIDATION RESISTANCE  AND METHODS FOR IMPROVEMENT  A. Oxidation  The  oxidation  performance  of  zirconium diboride 1960s,[6,7]  has  been  investigated  since  the  and  the  temperature  range  in which  the  ceramics  have  been  tested has  increased with the testing capabilities of  the  time.  For monolithic  ZrB2, which  the  oxide  formed  on  oxidation is B2O3, (450 °C) and wets the oxide grains until it volatilizes at temperatures (1100 °C). 1373 K (1100 °C), the oxidation kinetics are controlled by the diffusion of oxygen through the liquid boria that  is  liquid  above  723 K  above  1373 K  Below  surrounds the ZrO2 grains. 1473 K (1100 °C and 1400 °C), display paralinear characteristics because of  Between  1373 K  and  the oxidation kinetics  the mass  gain from ZrO2 and B2O3 from B2O3 vaporization. It has long been known  formation,  and mass  loss  that  the  addition  of  SiC  improves  the  properties  of  ZrB2 concluded generally  UHTCs  signifi cantly,[4,8-12]  and  it  is  that  the  optimum amount of SiC is between 15 and 20 vol pct.  The addition of SiC increases  the  sinterability of  the  starting powders by facilitating liquid phase sintering via  formation of a borosilicate liquid. This  in turn allows  greater control of the diboride grain growth which has a effect on mechanical properties.[13,14] Other  beneficial  benefits reported include improved thermal shock resisand oxidation resistance.[15] The  tance  impact of  the  presence of a connected low thermal conductivity and  brittle  glassy phase on thermal shock is likely to be below 1273 K (1000 °C). During oxidation at elevated temperatures, the silicon  limited  since  it will  soften  from the SiC and boron from the ZrB2 are oxidized to  EMILY EAKINS, PhD Student, DONI DANIEL JAYASEELAN,  Post Doctoral Research Associate and WILLIAM EDWARD LEE,  Professor,  are with  the Centre  for Advanced  Structural Ceramics  (CASC) and Department of Materials, Imperial College, London SW7  2AZ, U.K. Contact e-mail: w.e.lee@imperial. ac.uk  Manuscript submitted June 7, 2010.  Article published online November 23, 2010  878—VOLUME 42A, APRIL 2011  METALLURGICAL AND MATERIALS TRANSACTIONS A  \\x0c', 'form a protective borosilicate glass  layer according to  the following equations:  ZrB2 ðsÞ þ 5 2 SiC(s) þ 3 2  O2 ðgÞ ! ZrO2 ðsÞ þ B2O3  l,g  ð  Þ  O2 ðgÞ ! SiO2 ðl) þ CO gð  Þ  Several researchers[10,16-18] found that  the oxide scale  generally has a layered structure (Figure 1 and Figure 2).  A borosilicate glass  layer  is present at  the  top of  the  micrograph. Below this lies an oxide layer whose pores  have become filled with a glassy phase. The high wetting  angle between the oxide  grain and borosilicate  glass  ensures  complete  protective  coverage  of  the  exposed  faces below the temperature at which the glass evapo rates. In some samples, there is a region below this oxide  layer  that  is depleted in SiC. Finally, at  the bottom of  the micrographs is the unreacted bulk ceramic.  The borosilicate glass has a higher viscosity, higher  boiling  point,  and  lower  vapor  pressure  than  boria,  providing more  efficient  oxidation  protection.  The  effectiveness  of  the  protective  layer  increases  up  to  approximately 20 vol pct SiC. When ZrB2-SiC materials undergo furnace oxidation at 1773 K (1500 °C), the layer decreases  thickness  of  the  borosilicate  with  increasing SiC content. When sufficient SiC is present,  the borosilicate liquid flows  from the site of oxidation  toward the surface of the material and as the amount of  SiC increases,  the liquid can fill  the spaces between the  oxide  grains  efficiently  to  protect  the material  from  subsequent oxidation.  The  formation and efficacy of  the oxidation micro structure has undergone extensive investigation. et al.[19] provided a detailed explanation of  Li  the mecha nisms  responsible for ZrB2-SiC UHTCs dation performance over both monolithic ZrB2 and SiC. They suggested that the improved oxidation perfor improved oxi mance is caused by the development of a ‘‘solid pillars  liquid roof’’ structure in which the borosilicate glass acts  as a diffusion barrier and the ZrO2 grains liquid and provide mechanical stability to  retain the  the  oxide  layer. The  oriented  growth  of  the ZrO2 ceramics is caused  grains  on  oxidation  of  ZrB2-SiC production of gaseous by-products and direction of liquid transport.[20] The removal of ZrO2 grains occurs by the discharge of SiC gaseous products with high  by  the  vapor pressures:  carbon monoxide gas during passive  oxidation and carbon monoxide and silicon monoxide  gases during active oxidation. ZrO2 can be transported to the surface of the glass layer by convection or by reactions between the zirconia and the boria or silica.[20]  Tailoring  the  glass  composition  to  inhibit  oxygen  transport could slow down the progression of the oxide  layer into the underlying bulk.  It  is believed that  the borosilicate liquid moves from  the site of oxidation to the surface. This is thought to be  caused by the silica liquids being nonwetting for zirconia  and by the liquid phase being more viscous nearer  the  surface because of B2O3 volatization. Karlsdottir and Halloran[21] proposed that continued recession of the  oxidation subscale during oxidation is  caused by  the  flow of boria-rich liquid in ‘‘convection cells’’ between  the oxide grains that provide transport paths for oxygen  to progress  into the, as yet unoxidized, bulk material.  Fig. 1—Scanning  electron microscope  (SEM) image of ZrB2 + 20 1900 K (1627 °C) for ten 10-min cycles. The numbers on the micrograph indicate: (1) outer  vol pct  SiC,  after  oxidation  in  air  at  glass layer, (2) ZrO2-based interlayer, (4) unreacted bulk ceramic.[10]  (3)  SiC depleted zone,  and  Fig. 2—SEM micrograph of  the oxide  scale  in a ZrB2 + 20 vol pct SiC + 5 vol pct Si3N4, after nonisothermal oxidation testing up to 1623 K (1350 °C). The numbers on the micrograph indicate: (1) outer glass layer, (2) ZrO2 -based interlayer, and (3) unreacted bulk.[16]  METALLURGICAL AND MATERIALS TRANSACTIONS A  VOLUME 42A, APRIL 2011—879  \\x0c', 'The convection cells within the borosilicate liquid result  in precipitation of zirconia on the oxidized surface as is  shown in Figure 3 and demonstrated schematically in  Figure 4.  The amount of SiC must be controlled so that there is  sufficient protective glass  formed on oxidation. How ever,  if  the  amount  of  SiC  present  is  above  the  percolation threshold,  i.e.,  it  is present as an intercon nected three-dimensional network,  then open channels  of porosity remain between the oxide grains when it  is  oxidized. These will then leave the remaining SiC subject  to additional oxidation and lead to a SiC depletion zone. Peng et al.[22] argued that  the SiC depleted zone is  the  result  of  a wicking  process.  The  internally  formed  borosilicate liquid is transported via capillary action to  the existing borosilicate liquid on the surface formed by  earlier oxidation. Increasing the SiC content above the  percolation threshold increases the thickness of the SiC depleted zone.[20] The SiC depletion zone has also been  observed to act as  the  initiation site  for  cracking and  spallation of material (1900 °C). ZrB2 grains adjacent to zones are left exposed to oxygen.[20]  during  oxidation  at  2173 K  the  SiC-depleted  The size of the SiC powder particles also has an effect  on the properties of  the UHTCs. The use of ultrafine  SiC with submicron particle size can produce ZrB2-SiC composites that have improved mechanical and ther momechanical  properties  compared with  a material  prepared with standard starting powders of much larger size.[23] This  particle  is  because  of  a more  uniform  particle dispersion that suppresses diboride grain growth  resulting in a finer microstructure that  is more efficient  at  crack  deflection  and  bridging. Ultrafine  SiC also  allows material to be produced 2173 K (1900 °C) without and improves the oxidation material.[24-27] The flexural  by  hot  pressing  at  the presence of a sintering  aid[24]  resistance  of  the  strength of ZrB2-SiC composites produced with ultrafine SiC can increase after  oxidation, whereas  the flexural  strength of composites  produced with larger SiC powder (average particle size \\x196.4 lm) decreases after oxidation.[26] As the temperature is increased, a passive  to active  transition of SiC oxidation occurs between 1873 K and 1973 K (1600 °C and 1700 °C) in air under atmospheric pressure,[25,28] so that above these temperatures, any  protection afforded by a glassy oxide layer is largely lost.  III.  OXIDATION KINETICS  Extensive oxidation studies have been performed to  determine the oxidation kinetics of ZrB2 and ZrB2-SiC composites.[4] For the oxidation of ZrB2 without SiC, the diffusion of oxygen through the boria glass layer was identified as the rate limiting step up to ~1473 K (1200 °C). The protective liquid phase evaporates above this temperature; thus, the oxidation rate becomes  dependent on oxygen diffusion through the  zirconium  oxide layer.  UHTCs  containing SiC retain the protective glassy  layer for a larger temperature range than those without.  However, SiC only improves the oxidation resistance at temperatures above 1623 K (1350 °C), at which the silicon is oxidized. Below this temperature,  the temperature  Fig. 3—Backscattered electron SEM images of (1600 °C).[21]  convection cells on the surface of ZrB2 + 15 vol pct SiC after oxidation for 30 min at 1873 K  Fig. 4—Schematic of convection cell  features on the surface of oxi dized ZrB2-SiC UHTCs. BSZ is a liquid oxide boria, silica, and zirconia.[21]  solution containing  880—VOLUME 42A, APRIL 2011  METALLURGICAL AND MATERIALS TRANSACTIONS A  \\x0c', 'SiC inclusions remain in the oxide layer. Hinze et al.[29]  found  that  the  rate-limiting  step  accounting  for  the  parabolic kinetics observed in these materials was  the  diffusion of oxygen through the borosilicate glass; this idea has been supported by more recent research.[13] Bellosi[13]  Monteverde  and  studied  the  oxidation  resistance of hot-pressed HfB2-SiC composites by isothermal and nonisothermal treatments in air at temper£1873 K (1600 °C). oxidation kinetics of the composite fit a paralinear law  atures  They  found  that  the  until  partial  rupture  of  the  oxide  scale  occurs,  after  which the weight gain fits a linear  law. The parabolic  contribution is  a  result of  the growth of an external  oxide  scale,  which  imposes  longer  diffusion  paths  progressively on oxygen arriving at  the diboride-oxide  interface. The negative linear contribution accounts for  volatile  products  (most  likely B2O3 outermost borosilicate  or  SiO2) layer. The  being  released  from the  HfB2-SiC composite also contained HfO2 and a cubic Hf(C, N) solid solution as secondary phases formed  during hot pressing. These had a detrimental effect on  the oxidation resistance at temperatures below 1623 K (1350 °C). At temperatures above 1623 K (1350 °C), the SiC improved the oxidation resistance as expected.  While investigating the HfB2-SiC material, Monteverde and Bellosi[13] suggested the main reactions that  described  the  oxidation  process. Depending  on  the  selected  temperature  range,  these  reactions  involve  either mass gain, as follows:  HfB2 ðsÞ þ 5 2 O2 ðgÞ ¼ SiO2 ðsÞ þ CO(g)  O2 ðgÞ ¼ HfO2 ðsÞ þ B2O3 ðl)  SiC(s) þ 3 2 2BN(s) þ 3 O2 ðgÞ ¼ B2O3 ðl) þ N2 ðg) HfC1\\x00xNx ðsÞ þ 3 \\x00 x 2 2 ¼ HfO2 ðsÞ þ 1 \\x00 x  O2 ðgÞ  ð  ÞCO(g) þ x 2  N2 ðgÞ  or mass loss, as follows:  B2O3 ðl) ¼ B2O3 ðg) SiO2 ðs) þ CO(g) ¼ SiO(g) þ CO2 ðg) 2SiO2 ðs) þ SiC(g) ¼ 3SiO(g) þ CO(g)  For  a UHTC material  to maintain  its  oxidation  resistance,  it must retain the protective oxide layer. and Bellosi[23] observed rupturing of  Monteverde  the  oxide scale in a ZrB2-SiC ceramic after tion testing through several thermal cycles 1973 K (1700 °C). The the oxide scale was attributed in part  furnace oxida from room  temperature  to  cracking  and  spalling of  to a  phase  transformation in the ZrO2, which takes place during thermal cycling. At high temperatures, the oxide  exists in tetragonal form, whereas at lower temperatures,  it  exists  in monoclinic  form.  For  pure  ZrO2, athermal martensitic monoclinic to tetragonal transfor1443 K (1170 °C) on on cooling.[30,31] This  the  mation occurs at 1223 K (950 °C)  heating  and  is  obviously  undesirable for a material that will be subject to thermal  cycling  with  rapid  thermal  transient  conditions.  The  oxides  also  have  higher  coefficients  of  thermal  expansion (CTEs) and lower thermal conductivities than  the underlying diboride material. This, coupled with the  phase  transformation,  leads  to cracking of  the oxide  scale, allowing oxidation of the underlying bulk.  At  low partial  pressures  that  exist  close  to  the  interface  between  the  unreacted  bulk  and  the  oxide  layer, ZrO2 forms oxygen lattice vacancies and becomes nonstoichiometric (i.e., ZrO2-x). This increases the through the  diffusivity  of  oxygen  oxide  scale, which  also increases oxidation of the underlying bulk.  IV.  ADDITIVES  The oxidation products of a material  formed during  exposure  to  an  oxidizing  environment  are  largely  responsible  for  the  high-temperature  performance  of  that material. This is determined by the extent to which  the oxidized layer  can protect  the bulk material  from  subsequent oxidation. The physical and chemical pro cesses that occur at  the exposed surface depend on the  microstructure and composition of  the oxidized mate rial.  It  follows  that modification of  the microstructure  and composition can have a beneficial  (or detrimental)  effect on the material’s oxidation resistance.  Additives can be used in several ways to improve the  oxidation resistance of  the UHTCs. The main areas of  interest are as follows:  (a)  Increasing viscosity of the borosilicate liquid  (b)  Inhibiting the ZrO2 polymorphic transformations (c) Using alternatives to SiC to introduce silicon  (d) Forming protective refractory phases at high tem perature  (e) Modifying the microstructure of the ZrO2 scale  A.  Increasing Viscosity of  the Borosilicate Liquid  Systems with higher viscosity and increased liquidus  temperatures  inhibit oxygen diffusion to the unreacted  bulk, retain the protective liquid at higher temperatures,  and  suppress  evaporation  of  boria  from the  glassy  phase. Diffusivity  is  inversely  proportional  to  the  viscosity of  the  liquid through which the diffusion is  taking place, which relationship[32]  is  shown  in  the  Stokes-Einstein  D ¼ kT  6pgr  where D is the diffusion constant, k is Boltzmann’s constant, T is temperature in Kelvin, g is viscosity, and  r is the spherical particle radius.  The viscosity of the borosilicate glass can be increased  even more  by  adding  certain  elements  to  the  bulk  material. The addition of  tungsten in 10 and 20 vol pct  additions to ZrB2-SiC ceramics increased the viscosity of the borosilicate glassy phase significantly but also  reduced  the material’s  thermal  shock  resistance  and  structural  stability  at  elevated temperatures, which is leading-edge materials.[4]  highly undesirable for sharp,  METALLURGICAL AND MATERIALS TRANSACTIONS A  VOLUME 42A, APRIL 2011—881  \\x0c', 'The  oxidation  resistance  of  hot-pressed ZrB2 + 25 vol pct SiC composites has been improved by the addition of diborides of Cr, Ti, Ta, Nb, and V.[33] These  additions  result  in  the  production  of  the  respective  oxides  in  the  glass. The  improvement  in  oxidation  resistance comes  from the fact  that borate and silicate  glasses containing oxides of  the elements listed (Group  IV-VI  transition metals)  are  immiscible  and  lead  to  phase  separation. Such systems  contain compositions temperatures.[32] The  with high viscosity and liquidus  immiscibility  of  the  glass  increases with  increasing  cation  field strength z/r2 where  of  the metallic  oxide  forming  element  z  is  the valence and r  is  the  ionic  radius. ZrB2 CrB2, NbB2, TaB2, oxidation resistance,  ceramics with  additions  of  10 mol pct  TiB2, but  and  VB2 improvement  had  improved  most  was  observed in the ZrB2-SiC + 10 mol pct TaB2, which displayed a weight gain of ~1.25 pct after 5 hours of thermogravimetric analysis (TGA) compared with ~3.5 pct for the ZrB2-SiC material with no additives.[33] The order in which the oxidation resistance was  improved correlated well with the cation field strength  of  the modifying additive. et al.[11]  Opila  investigated  the  effect  of  tantalum  additions on the oxidation performance of  zirconium  diboride.  They  found  that  the  addition  of  TaSi2 ZrB2 + 20 vol pct SiC composite. The oxidation rate was reduced by a factor of 10 at 1900 K (1627 °C). They concluded that more research was required to confirm that the  improved  the  oxidation  resistance  of  a  improvement  in oxidation was a result of  the tantalum  addition and not  from the  accompanying  increase  in  silicon.  It was  suggested that  the  introduction of Ta  resulted  in  immiscibility  of  the  liquid  formed  on  oxidation, which increased the viscosity of  the  liquid  phase,  providing  a  protective  layer  that was more  resistant  to volatilization. Evidence of glass immiscibil ity  on  the  surface  of ZrB2 after oxidation in air  20 vol pct SiC-20 vol pct TaSi2 1900 K (1627 °C) 100 minutes is shown in Figure 5.  at  for  Peng and Speyer[34]  investigated the  effect of TaSi2 and TaB2 on the oxidation resistance of ZrB2-B4C-SiC composites across the temperature range 1423 K to (1150 °C 1550 °C). Ta-containing compounds improved the oxidation resis 1823 K  to  The  addition  of  tance of the materials over the entire temperature range  studied, but  the TaSi2 performed better than the TaB2 presumably because of the ability to form larger amounts of protective silica-based liquid. Peng et al.[22]  investigated the effect of SiC, TaB2, and TaSi2 on the isothermal oxidation resistance of ZrB2. They found that an increase in SiC decreased the thickness of both  the silica-rich glassy layer and the SiC depletion zone.  TaB2 was more performance than TaSi2. Both to be a result of the formation of a zirconium effective  at  improving  the oxidation  caused  improvements  thought  tantalum boride  solid solution. After oxidation of  the  solid solution, segregated ZrO2 and TaC phases resulting in finer particles (~1 lm) present in the liquid phase than with oxidation of pure ZrB2. The trapping the liquid phase  form,  finer  particles are more effective at  in  the ZrO2 through the liquid. The additives were effective only at (~3.32 mol pct) concentrations were detrimental to the oxidation resis layer  and  preventing  oxygen  transport  small  concentrations  and  at  larger  tance because of  the  formation of  zirconia dendrites,  which act as conduits for oxygen transport into the bulk.  The addition of yttria has been investigated by several et al.[35]  groups. Zhang  found  that  adding  3 vol pct  improved sinterability of  the powders and suppressed  grain growth by  reacting with oxides on the  starting  powder  surfaces. Grain size  refinement  improved the  fracture toughness and flexural strength of the material. SiC UHTC[36]  Adding LaB6 a ZrB2 + 20 vol pct resulted in significantly higher fracture toughness com to  pared with the same UHTC without the LaB6 (5.7 MPa m1/2 and 4.0 to 4.8 MPa m1/2, respectively) because of  enhanced crack deflection and bridging near SiC parti cles. MoSi2 has a beneficial effect on the mechanical and oxidation properties of ZrB2-SiC ceramics and is an effective sintering aid for ZrB2-SiC UHTCs produced by hot pressing[15,37] and spark plasma sintering.[38] Zhang et al.[37]  found  that  additions  of  tungsten  carbide  improved the oxidation resistance of pressureless 1873 K (1600 °C) the ZrB2, which allowed the ZrO2 formed during oxidation. This process resulted in a substantial decrease in oxide scale  sin tered ZrB2 solution with  at  by  forming  a  solid  liquid  phase  sintering of  thickness in the WC-containing ceramic.  B.  Inhibiting the ZrO2 Polymorphic Transformations  The integrity of  the oxide scale can be improved by  inhibiting the ZrO2 polymorphic their associated volume changes.  transformations and  In low temperature  systems,  this  is achieved by the addition of  stabilizing  cations such as Mg, Ca, and Y. However, these cations  are lost  from the ZrO2 lattice at relatively low temperatures, and for UHTCs, alternative cations have been  sought. The addition of a cation such as Ta results  in  substitution of  the cation on the Zr  site in ZrO2, reducing the concentration of oxygen vacancies because  thus  Fig. 5—SEM backscattered image showing evidence of glass  immis cibility on the surface of ZrB2 -20 vol pct SiC-20 vol pct TaSi2 after oxidation in air at 1900 K (1627 °C) for 100 min. Bright contrast  phase is ZrO2, intermediate contrast impurities and dark phase is SiO2.[11]  phase  is  silicate  glass with  882—VOLUME 42A, APRIL 2011  METALLURGICAL AND MATERIALS TRANSACTIONS A  \\x0c', 'of  the higher valence of  the  cation (Ta forms Ta2O5 when oxidized). This decreases oxygen diffusion through  the  scale  and  stabilizes  the  oxide  phase,  increasing  adhesion of the scale to the bulk ZrB2-SiC material. The cation must be of higher valence and must form a  refractory oxide  scale.  In addition to this,  candidate  additives must be introduced as a refractory phase and  form a refractory oxide. The  two best  candidates are  those based on niobium and tantalum, but  tantalum is  preferable as Ta2O5 2153 K (1880 °C) (compared with for Nb2O5).[33] Tantalum can form or as a carbide, boride, or silicide. The formation  has  a melting temperature of 1793 K [1520 °C] added in elemental  be  of  intermediate  phases  should  be  considered.  For  instance,  the  addition  of Ta2O5 6ZrO2 with ZrO2. This phase has temperature than the pure  could  form Ta2O5. lower melting  a  oxides  and  could  have  a  beneficial or detrimental effect on the oxidation behav ior of the composite.  C. Using Alternatives to SiC to Introduce Silicon  Alternative methods of  introducing Si  to the system  have been investigated. Ceramics  in the  system ZrB2Ta5Si3 have been investigated, as Ta5Si3 has a higher than SiC (2773 K and 1573 K [2500 °C melting point and 2300 °C], respectively). Ta5Si3 provides the tantalum to induce glass immiscibility and silicon to form the layer.[39]  protective  borosilicate  glass  in  the  oxidized  ZrB2and HfB2-TaSi2 composites have been formed by hot pressing,[40] resulting in formation of a solid solution  as  the tantalum entered the boride matrix. The HfB2TaSi2 displayed superior mechanical properties to the ZrB2-based material. The substitution of Ti for (e.g., a Zr-B-Ti system) has been suggested for making  SiC  an alternative UHTC system because of the significantly  reduced vaporization rate displayed by TiO2 compared with SiO2.  [41]  D. Forming Protective Refractory Phases at High  Temperature  Research into the introduction of additives  to ZrB2largely focuses on using the additives to  SiC ceramics  alter  the  properties  of  the  liquid  phase  formed  at  oxidation. A different approach is  to use additives  to  form a solid refractory scale at high temperatures, which  can resist oxidation  at higher  temperatures  than the  original material,  thus providing effective protection to  the underlying bulk and preventing subsequent oxida tion. Candidate  additives  for  this  approach  include  those based on rare earth elements,  in particular those  containing lanthanum.  Originally, zirconium diboride was investigated as an  additive to improve the oxidation resistance of LaB6 and was found to be effective up to ~1573 K (1300 °C).[42] Zhang et al.[43] prepared hot-pressed ZrB2-20 vol pct SiC-10 vol pct LaB6 and compared the oxidation performance with a ZrB2-20 vol pct SiC ceramic by oxidizing both materials with an oxyacetylene torch. Both samples underwent oxidation up to 2673 K (2400 °C), and the ceramic with the LaB6 addition displayed significantly less spalling and cracking than the ZrB2-SiC sample the weight changes for the ZrB2-SiCLaB6 and ZrB2-SiC were 0.2 pct and 3.1 pct, respectively. An energy-dispersive X-ray analysis of the oxidized  (Figure 6). Also,  surfaces  showed that  the LaB6-containing sample had formed an oxidized layer containing m-ZrO2, t-ZrO2, La2O3, and La2Zr2O7. The addition of LaB6 not only impeded the tetragonal to monoclinic ZrO2 transformation on cooling but also formed a self-generating  oxidation  barrier  containing  lanthanum  zirconate,  which  has  a  significantly  higher melting  temperature  than SiC. The elemental maps  showed that both silica  and boron are no longer present in the outer oxide layer after oxidation at ~2673 K (2400 °C), but num additions have been retained in a compact scale on  the  lantha the oxidized surface.  Fig. 6—Photographs of (2400 °C) for 600 s.[43]  (a) ZrB2-20 vol pct  SiC and  (b) ZrB2-20  vol pct  SiC-10 vol pct LaB6  after  oxyacetylene  torch  testing  at  2673 K  METALLURGICAL AND MATERIALS TRANSACTIONS A  VOLUME 42A, APRIL 2011—883  \\x0c', 'Lanthanum has also material as La2O3,[44] but an amorphous grain boundary phase  been  added  to  a  ZrB2-SiC resulted in the formation of  and substantial  ZrB2 and SiC grain growth. The same work found that additions of other rare earth oxides (Y2O3 and Yb2O3) had beneficial effects on the densification, hardness, and  fracture toughness of  the ZrB2-SiC but did not tigate their effect on oxidation resistance. Jayaseelan[45]  inves investigated the addition of several rare  earth (RE)-containing compounds to ZrB2 + 20 vol pct SiC. Samples were prepared with 10 vol pct LaB6, La2O3, or Gd2O3 and underwent oxidation testing at 1873 K (1600 °C). All samples successfully formed a (>100 lm), thick dense of RE2Zr2O7 oxidation (Figure 7). These zirconates have melting temperatures >2573 K (2300 °C) and will provide oxidation protection at temperatures when the borosilicate  layer  during  phase has vaporized from the exposed surface. Also, the  reaction of the RE with ZrO2 is expansive and therefore fills voids at the oxidized surface created by the removal  of volatile species such as B2O3.  E. Modifying the Microstructure of the ZrO2 Scale  Another novel  technique  is  the use of  additives  to  alter the microstructure of the ZrO2 scale. By providing a liquid phase sintering route for the ZrO2, it is possible to decrease the porosity of the scale and inhibit  the  subsequent transport of oxygen into the bulk material. As ZrO2 has a melting point of 2988 K (2715 °C), a sufficiently dense scale would provide effective oxidation  resistance  at  temperatures  above  those  at which the (~1873 K investigated by  boria or borosilicate [1600 °C]). This approach Zhang et al.[46] with additions of W to ZrB2, which results in formation of a WO3-ZrO2 eutectic at ~1548 K (1275 °C). ZrB2 + 4 mol pct WC ceramics underwent TGA at 10 deg/min to 1773 K (1500 °C) and isothermal 1873 K (1500 °C or oxidation studies at 1773 K or 1600 °C) in flowing air. The ZrB2 + 4 mol pct WC had better oxidation resistance than ZrB2, as indicated by the normalized mass gain after TGA heating to 1773 K (1500 °C) (~4.5 mg/cm2 and ~14 mg/cm2, respectively). ZrB2 + WC also showed superior oxidation resistance in the isothermal oxidation tests at 1773 K and 1873 K (1500 °C and 1600 °C), and the reduced mass gain of the ZrB2 + WC samples was longer oxidation times, supporting  phase  is  vaporized  has  been  for  1,  2,  or  3 hours  more  significant at  densification of  the ZrO2 scale. The addition of WC to ZrB2-SiC ceramics was also investigated and the presence of W increased the oxidation resistance of these  ceramics as well. However, additional  research is nec essary to confirm that the improved oxidation resistance  is caused by liquid phase sintering of  the ZrO2 scale.  V.  DENSIFICATION METHODS  The conventional method for densifying UHTCs is hot  pressing,  performed  either with  or without  sintering  additives. Significant  research has  examined alternative  processing routes for UHTCs to reduce processing times  and temperatures and, therefore, reduce the cost associated  with the techniques. Some of these alternative techniques  also improve the oxidation resistance of the material.  The presence of secondary phases in the microstruc ture  has  a  detrimental  effect  on  the material  high temperature  capability by introducing grain boundary  phases, which can have lower melting temperatures and  provide  routes  for oxygen diffusion into the material.  The presence of oxygen impurities can be detrimental to  the densification ability of the starting powder or cause  rapid grain growth;  it also contributes to the formation  of  secondary phases. Nitrides  and reducing  additives  have been added to the powders to enhance sinterabil ity,[23,47-49]  but  these  additives  introduce  secondary  phases to the material.  The main production techniques used for  formation  of UHTC materials are as follows:  (a) Hot pressing  (b) Pressureless sintering  (c) Self-propagating high-temperature synthesis (SHS)  (d) Reactive hot pressing (RHP)  (e) Spark-plasma sintering (SPS)  Fig. 7—Secondary  electron  micrographs  of  oxidized  scale  on  (a) ZrB2-20 vol pct SiC-10 vol pct LaB6 showing (1) dense La-containing layer, containing layer, and (3) unreacted bulk.[45]  SiC-10 vol pct La2O3 after oxidation for  and  (b) ZrB2-20 vol pct 1873 K (1600 °C) at  1 h  (2)  intermediate  ZrO2 884—VOLUME 42A, APRIL 2011  METALLURGICAL AND MATERIALS TRANSACTIONS A  \\x0c', 'A. Hot Pressing  Hot  pressing  is  the  conventional  method  for  fabricating  UHTCs  and  has  been  used  extensively,[8-10,12,16,17,20,22,27,33,35,39,47,48,50-65] with typi cal  temperatures  of  ~2173 K (1900 °C) 30 and 50 MPa. It  and  applied  pressures  between  allows  full  densification without the use of sintering aids, although  most research employs modest amounts of sintering aids  such  as  silicides,  borides, metals  (e.g., Ni),  or C to  reduce processing times and temperatures, thus reducing  the costs associated with the production technique. et al.[66]  Monteverde  have  performed  extensive  research using hot pressing and various  sintering aids.  A monolithic ZrB2 ceramic was compared with ZrB2TiB2 and ZrB2-B4C composites, and it was concluded that the composite materials had better mechanical  properties and performed better  in long-term furnace  oxidation  tests.  Full  densification  of  a HfB2 + 30 vol pct SiC ceramic was obtained by hot pressing at 2173 K (1900 °C) for 35 minutes using 2 vol pct TaSi2 as a sintering aid.[67]  B. Pressureless Sintering  The  refractory nature of ZrB2 makes sinter, and in general, pressureless sintering requires the  it difficult  to  presence  of more  sintering  aids  than  other  firing  techniques. As  a  consequence,  pressureless  sintered  material often contains amorphous phases  that can be  detrimental to the high-temperature capability of the material. Chamberlain et al.[68] produced a ZrB2 ceramic without sintering aids that reached 98 pct theoret ical density but required 9 hours pressureless sintering at 2423 K (2150°C). Fahrenholtz et al.[69] found that ZrB2 could not be sintered pressurelessly without the presence  of a sintering aid (in this case a combination of B4C and C), which reacts with and removes oxygen impurities the diboride powders. A ~100 pct formed at 2173 K (1900 °C) with B4C and C present as sintering aids.  from the surface of  dense material was  C. Self-Propagating High Temperature Synthesis  SHS is not a densification method but uses solid-state  combustion  to  produce materials  by  using  internally  generated chemical  energy from exothermic  reactions.  The characteristics of  this method include fast reaction  times,  low energy  requirements,  simple  experimental  apparatus, and high-purity products. A disadvantage is  that  the reactions are difficult  to control.  SHS can also be used to prepare ZrB2 powders using inexpensive raw materials.[70] When used for the pro duction of powders,  the high heating and cooling rates  involved are thought to introduce planar defects, such as  stacking faults, and linear defects  such as dislocations  whose associated strain fields increase the sinterability of  the powders by providing a driving force for rearrange ment of atoms. ZrB2 powders can be formed by SHS using zirconium and boron.[71] The powders are dry  mixed and cold pressed to form pellets. The pellets are  ignited and the process is an explosive one accompanied  by a large gas  release. As a consequence,  the resulting  material  is  too fragile  to be used as a bulk material,  although X-ray diffraction indicates it  is ZrB2. conventional hot  SHS has  several  advantages over  pressing  of UHTCs,  such  as  control  of  exaggerated  grain growth and lower processing temperatures. How ever,  several problems with hot pressed materials have  been overcome by addition of  reinforcing phases and  refinement of starting powder size.  D. Reactive Hot Pressing  To avoid the expensive processing conditions of hot  pressing, RHP can be used as an alternative production  route. RHP involves in situ high-temperature, solid-state  chemical displacement reactions. This has the benefit of  controlling the microstructure and producing chemically  compatible, evenly distributed phases. The main advan tage of RHP compared with SHS from a processing  point of view is that the displacement reactions can take  place at much lower temperatures than those that occur  during  SHS.  The  heating  rate  during reactive (~10 °C/min)  hot  pressing must  be  sufficiently  slow  to  prevent spontaneous self-combustion. Monteverde[72]  fabricated a HfB2-SiC composite by reactive hot pressing. Solid reagents (Hf/Si/B4C) were converted into the basic compo mixed mechanically,  nents (HfB2 and SiC), and hot pressed directly at 2173 K (1900 °C) to achieve full density. The composite exhibited comparable physical properties to conventionally hot pressed material. Wu et al.[14] produced a ZrB2SiC-ZrC composite from a mixture of zirconium,  silicon, and B4C by (1800 °C) for 60 minutes under atmosphere. The microstructure was not homogeneous;  reactive hot pressing  at  2073 K  20 MPa  in an argon  it contained residual porosity and large ZrB2 grains (up to 10-lm diameter) because of the relatively large size of the starting powder (£25-lm particle size). These problems could be rectified by improving the uniformity of  mixing and using starting powders with smaller particle  size. The  starting powder  size and morphology has a  strong effect on the microstructure et al.[73]  in RHP ceramics.  Qiang  formed  a  fully dense ZrB2-SiC-ZrC composite successfully at 2173 K (1900 °C) by RHP. The microstructure was fine grained (~1 lm), resulting in high values of flexural strength (526 ± 9 MPa) and fracture toughness (6.50 ± 0.30 MPa m1/2). Wu et al.[74] fabricated a ZrB2-SiC-ZrC composite by RHP at 2073 K (1800 °C)[14] and another at 1873 K (1600 °C), which reached theoretical densities of 96.8 pct and  97.3 pct, respectively. The material sintered at 1873 K (1600°C) displayed mechanical properties comparable to those previously reported for materials that had expe rienced  higher  sintering  temperatures.  Rangaraj  et al.[75,76]  found that nonstoichiometric ZrCx played a crucial role in RHP of Zr-B4C powder mixtures at 1473 K (1200 °C) and in ZrB2-SiC composites produced at 1873 K (1600 °C). Zhang and Zhang[62] compared the effect of process ing method on microstructure by examining ZrB2-SiCZrC composites that had been formed by RHP and hot  METALLURGICAL AND MATERIALS TRANSACTIONS A  VOLUME 42A, APRIL 2011—885  \\x0c', 'pressing. The hot-pressed material had equiaxed ZrB2 grains, whereas the RHP material had equiaxed and  plate-like ZrB2 grains. This difference in microstructure had an effect on the mechanical properties of the  materials: The HP material had higher flexural strength  (681 ± 67 MPa compared with 654 ± 17 MPa for  the  RHP material), whereas  the RHP material  exhibited  more effective toughening mechanisms.  E. Spark-Plasma Sintering  SPS, which is also known as the field-assisted sinter ing technique or pulsed electrical current sintering,  is an  unconventional method for consolidating powders with  relatively  poor  sinterability  in  short  times.  SPS  uses  conventional  resistance heating and pressure but also  pulses a direct  current  through the graphite die  (and  powder  if  it  is electrically conducting), which is  shown  schematically in Figure 8. The pulsed current enhances  grain boundary diffusion and migration, which increases  the final density of  the  ceramic by eliminating closed  porosity.  SPS has been shown to provide a processing route for  these highly refractory materials  that does not  require  sintering  aids  and  produces  fully  dense  products  at  lower temperatures and in shorter tional hot pressing.[78,79] This has  times  than conven the  effect of main taining  a  fine  grain  size, which  is  beneficial  to  the  mechanical properties of  the material  such as  fracture  toughness. The technique has been particularly success ful when using starting powders fabricated using an SHS method.[79] researchers[78] have produced dense UHTCs  Several  successfully by SPS without sintering aids using sinter ing temperatures in the range of 2173 K to 2373 K (1900 °C to 2100 °C) and holding times on the order of minutes. Materials produced by SPS have displayed  comparable  and/or  improved mechanical  properties  and  oxidation  resistance when  compared with  hot pressed materials. The reason for the superior oxidation  resistance may be the removal of surface oxides during  sintering. The  storage  and  processing  of ZrB2 HfB2) in air results in the formation of surface oxides on the powder particles.[51] These surface oxides not only  (and  inhibit  the  sinterability of  the diboride, but also they  introduce oxides  such as B2O3 these low-melting-point oxides has a  into the bulk material.  The presence of  detrimental  effect on the high-temperature mechanical  properties and oxidation resistance. The pulsed electri cal current used during SPS is thought to cause thermal  or electrical decomposition of insulating surface oxides,[80,81] although the process by which the SPS method  achieves  this  and the  rapid sintering  rates  is not  yet  understood fully.  The parameters of the SPS technique can be altered to  produce materials with fine microstructures, which tend  to  contain  fewer  secondary  phases  than materials  produced by conventional methods. SPS materials can aids[82]  be  produced  using  sintering  and  still  have  superior oxidation resistance than hot-pressed materials.  VI.  CONCLUSIONS  There  are many methods  to  improve  oxidation  resistance  in UHTCs, but  they can be detrimental  to  other material properties or only provide improvements  over  limited  temperature  ranges. By  controlling  the  starting powder  size and additives, employing efficient  mixing techniques and using production methods  that  minimize  secondary phase  formation,  it  is possible  to  improve  the  oxidation  resistance.  Even  so,  current  research is  still  far  from producing a material  that  is  reliable and can be used repeatedly in extreme environ ments. 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},{
  "_id": 274,
  "PDF": "UHTC–carbon fibre composites Preparation, oxyacetylene torch testing and characterisation.pdf",
  "Text": "['Available  online  at  www.sciencedirect.com  Journal of the European Ceramic Society 33 (2013) 423-432  UHTC-carbon fibre composites: Preparation, oxyacetylene torch testing and characterisation  A. Paul a,∗  , S. Venugopal a , J.G.P. Binner a , B. Vaidhyanathan a , A.C.J. Heaton b , P.M. Brown b  a Department of Materials, Loughborough University, LE11 3TU, UK b DSTL, Porton Down, Salisbury SP4 0JQ, UK  Received 22 June 2012; received in revised form 13 August 2012; accepted 21 August 2012  Available online 13 September 2012  Abstract  Current generation carbon-carbon (C-C) and carbon-silicon carbide (C-SiC) materials are  limited  to service  temperatures below 1800 C and materials are sought  that can withstand higher  temperatures and ablative conditions for aerospace applications. One potential materials solution is carbon fibre-based composites with matrices composed of one or more ultra-high  temperature ceramics (UHTCs);  the  latter are  intended  to protect the carbon fibres at high temperatures whilst the former provides increased toughness and thermal shock resistance to the system as a whole. Carbon fibre-UHTC powder composites have been prepared via a slurry impregnation and pyrolysis route. Five different UHTC compositions have been used for  impregnation, viz. ZrB2 , ZrB2 -20 vol% SiC, ZrB2 -20 vol% SiC-10 vol% LaB6 , HfB2 and HfC. Their high-temperature oxidation resistance has been studied using a purpose built oxyacetylene torch test facility at temperatures above 2500 C and the results are compared with that of a C-C benchmark composite. © 2012 Elsevier Ltd. All rights reserved.        Keywords: Ultra-high temperature; Oxyacetylene torch testing; Oxidation  1.   Introduction     Refractory  transition metal  borides  and  carbides  have extremely high melting points of over 3000 C and hence are referred  to as ultra-high  temperature ceramics (UHTCs). Even though  they have been studied since  the 1960s,  there has been recent interest in these materials as potential candidates for thermal protection systems on hypersonic vehicles. A developmental history of UHTC materials can be found in Opeka et al.1 and they have recently been reviewed by Paul et al.2 The initial selection of UHTC materials was based on  their melting  temperatures, however oxidation  temperature and  the melting points of  their oxides is, in fact, more critical. There are a number of materials with melting points over 3000 C, whose oxides also have melting points  in excess of 2500 C,  for example ZrB2 , HfB2 and HfC. These materials are widely studied for high  temperature       ∗  Corresponding author at: Department of Materials, Loughborough Univer sity, Loughborough, Leicestershire LE11 3TU, UK. Tel.: +44 1509223154;  fax: +44 1509223949.  E-mail address: A.Paul@lboro.ac.uk (A. Paul).  0955-2219/$ - see front matter © 2012 Elsevier Ltd. All rights reserved.  http://dx.doi.org/10.1016/j.jeurceramsoc.2012.08.018     applications as monolithic components, however single phase ceramics are significantly limited in this area as a result of their very poor  thermal shock and oxidation resistance.3 Even with the addition of a second or  third ceramic phase such as SiC or LaB6 , these materials do not possess the high temperature resistance, thermal shock resistance or fracture toughness because of the volatilisation and decomposition of the oxidation products.4 The desired properties may require the development of fibre reinforced UHTC composites to enable viable application development beyond 2500 C. For such hybrid materials, carbon fibre is  the preferred choice owing  to  its high strength, ready availability and ability to be formed into complex shapes,5 provided it can be protected from oxidation. There are a number  reports  in  the  literature describing  the preparation of fibre  reinforced composites  for UHT applications, of which those developed by Levine et al.6,7 are amongst the earliest. They studied the high temperature oxidation resistance of SiC fibre reinforced ZrB2-20 vol% SiC, prepared via filament winding, slurry impregnation and hot pressing, against that of non-reinforced ZrB2-20 vol% SiC at up  to 1927 C for periods of up  to 100 min.6 Whilst  the non-reinforced material showed the best oxidation protection at 1327 C and 1627 C, at                 \\x0c', '424   A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432           1927 C both compositions underwent severe degradation and bloating and  the authors expressed concerns about  the  thermal  shock  resistance of  the non-reinforced materials  in high heat flux, aeroconvective environments. The same group7 prepared UHTC composites using Zoltek Panex® 30 carbon fabric, allylhydridopolycarbosilane preceramic polymer, HfB2 and SiC powders, to create a graded structure, from a HfB2 -rich surface through  to a SiC-rich surface with Si-O-C pre-ceramic polymer throughout, although micro-cracks were present in the final composites. Oxidation  testing was carried out  in a  furnace at 1617 C and using an oxyacetylene flame at 1805-2015 C. Following cyclic heating in the furnace, a non-uniform HfSiO4 and monoclinic HfO2 surface was formed on the HfB2 -rich surface and a glassy SiO2 layer was  formed on  the SiC-rich surface. Damage to the carbon fibres in the furnace testing was found to be lower at the HfB2 -rich surface compared to the SiC-rich surface. In comparison, during  the oxyacetylene flame  testing  the HfB2 -rich surface suffered a greater degree of damage during a 4-min test than that experienced by the SiC-rich surface. These highlights the significant differences between the results arising from different  test methods  in which not only  temperature but also gas flow rates differ and the effect this has upon the surface reactions and damage. As a result, direct comparisons between different high temperature test methods are rarely meaningful. Tang et al.8 prepared a  range of UHTC composites using 2D carbon fibre (ex-PAN Toray T700) preforms and five different mixes of aqueous UHTC powder slurries based on ZrB2 , SiC, HfC and TaC. A pressure assisted  technique was used to  impregnate  the powders  into  the fibre preforms and  then pyrolytic carbon deposition was used  to hold  the powders  in place. Analysis showed the UHTC powders to be concentrated 2 mm deep. The hybrid UHTC in a surface layer no more than  composites were  tested using an oxyacetylene flame; different gas ratios were used  to obtain different  temperatures and heat −2 the compositions containfluxes. At 1800 C and 2380 kW m ing SiC demonstrated  the  lowest erosion depth. However, at the more aggressive conditions of 2700 C and 3920 kW m a C/C-ZrB2 composite outperformed  the other compositions. Recently Zhao et al.9 prepared 3D Cf-ZrC composites using a precursor impregnation and pyrolysis route, studied the mechanical properties and evaluated  the high  temperature  resistance using an oxyacetylene  torch. They reported  that  the formation of ZrO2 melt on the surface contributed to a superior high temperature performance. The present study is designed to investigate further the potential of  carbon fiber based UHTC  composites  for ultra-high temperature applications. A number of UHTC powder compositions were used to prepare composites and the high temperature performance was evaluated using a custom built oxyacetylene torch test facility. The composites were characterised before and after high temperature testing.  −2        2. Experimental  ZrB2 (Grade B, 1.5-3  \\u242em), SiC  (Grade UF-25, 0.45  \\u242em), LaB6 (Grade C,  2-3  \\u242em), HfB2 (325 mesh, <44  \\u242em)  and HfC  (325 mesh, <44 \\u242em) were procured  from H. C. Starck     ×  (H. C.  Starck GmbH, Goslar, Germany).  Prior  to  further processing,  they were  characterised using XRD  (Bruker D8 diffractometer, Bruker AXS GmbH, Karlsruhe, Germany), FEGSEM  (Leo  1530VP  FEGSEM, LEO Elektronenskopie GmbH, Oberkochen, Germany), EDS  (EDAX, EDAX  Inc., NJ, USA), XPS  (ESCALAB 5, VG Scientific, West Sussex, UK), BET  (Tristar 3000, Micromeritics  Instrument Corporation, Norcross, USA) and particle  size analysis  (Mastersizer 2000, Malvern Instruments Ltd., Worcestershire, UK). 30 mm dia.   17 mm  thick 2.5D needled Cf preforms with 23 vol% fibres were obtained  from Surface Transforms Plc., Cheshire, UK, whilst  phenolic  resin  (Cellobond  J2027L) with  a  car45.5% (at 900 bon content of  C under an  inert atmosphere) was obtained from Hexion Specialty Chemicals, B. V., Rotterdam, The Netherlands. UHTC powder/phenolic  resin/acetone slurries were prepared by ball milling the ingredients in a plastic container using alumina milling media  for 48 h. A  typical slurry composition consisted of 40 g of UHTC powder, 20 g phenolic  resin  and 12.5 g  acetone. 5 different UHTC powder/compositions were used  to prepare  the  slurries  including ZrB2 , ZrB2-20 vol% SiC (ZS20), ZrB2-20 vol% SiC-10 vol% LaB6 (ZS20-1La), HfB2 and HfC. The Cf preforms obtained  from Surface Transforms were impregnated with the prepared slurries using a squeeze impregnation  technique where  the preforms were fully  immersed  in a beaker containing the slurry and squeezed manually repeatedly to achieve maximum slurry  intake. Four composites were prepared for each composition and the impregnated preforms were dried in an air oven at 75 C for 4 h followed by curing at 150 C for 2 h. This entire cycle was repeated 3 times to maximise the amount of UHTC powder within the composite. After the third impregnation and curing, the samples were pyrolysed at 900 C for 2 h using a tube furnace under flowing argon (99.998% pure) −1 . After pyrolusing a heating and cooling rate of 1.5 × 5 mm deep hole was drilled at  C min ysis, a 10 mm dia.  the bottom of the composites and further densification was achieved using chemical vapour  infiltration (CVI) of carbon at Surface Transforms using a commercial process. Benchmark carbon-carbon (C-C) composites were also prepared by CVI that had not undergone UHTC powder  impregnation. The change  in mass of  the samples was recorded after each stage in the preparation process and the bulk density of the composites was measured geometrically. Representative  composites were mounted  in  epoxy  resin, cross sectioned and polished using a semi-automatic polishing machine  (TegraPol-25, Struers Ltd., Solihull, UK) with  successively finer diamond polishing discs. The final polish used a 1  \\u242em diamond slurry and  the samples were analysed using SEM (Leo 1530VP FEGSEM, LEO Elektronenskopie GmbH, Oberkochen, Germany) to find the depth of impregnation of the powder  into  the carbon fibre preform. The powder distribution and the efficiency of powder mixing were evaluated using EDS mapping (EDAX, EDAX Inc., NJ, USA). After CVI,  the samples were analysed using micro-CT (Metris X-Tek 160Xi, X-Tek Systems Ltd., Hertfordshire, UK) to determine again the depth of  impregnation and  the distribution of UHTC powder within the Cf preform.              \\x0c', 'A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432   425  Table 1  Mass and density of UHTC powder impregnated composites.  Composite  Average mass of carbon  fibre preform (g)  Initial   After UHTC powder  impregnation and CVI  Final bulk density −3 ) (g cm  C-C   Cf-ZrB2 Cf-ZS20   Cf-ZS20-1La   Cf-HfB2 Cf-HfC   7.0   6.9   6.9   6.8   7.0   7.0   26.2   29.9   28.8   26.7   29.3   30.3   1.86   2.11   2.01   1.91   ± ± ± ± ±   0.01   0.01   0.03   0.03  1.93   0.03 2.07 ±  0.04  can be observed  that  these  layers are somewhat denser as  they are rich with UHTC powder, which is ideal for the potential to offer better oxidation and erosion resistance. The level of powder impregnation achieved with  the pressureless slurry  impregnation technique was better than that achieved by Tang et al.8 ; they reported a dense outer layer thickness of just 0.75 mm, even with a pressure assisted technique. This could be due to a difference in viscosity between  the slurries, differences  in UHTC powder particle size or differences  in porosity between  the preforms; none of these parameters were reported previously.8 The depth of  impregnation was further analysed on carbon CVI densified composites using micro-CT and a representative image  is shown  in Fig. 3. The 10 mm dia.   5 mm deep hole  ×  Fig. 1. Oxyacetylene torch test rig. (1) Back face thermocouple, (2) water cool ing, (3) sample holder, (4) sample, (5) sample guide, (6) protective   insulation,  (7) oxyacetylene   torch,   (8) neutral density filter,   (9)   thermal   imaging camera  and (10) 2 colour pyrometer.  The  high  temperature  oxidation  performance  of  the UHTC  composites was  studied  along with  the  benchmark carbon-carbon composites utilizing a custom built oxyacetylene  torch  test rig, shown  in Fig. 1, with an oxygen rich flame (1:1.35 acetylene to oxygen ratio). The specimens were fixed in a water cooled graphite sample holder with three graphite bolts and a K  type  thermocouple, connected  to a data  logger, was placed in contact with the back face of the sample through a hole drilled in the sample holder to record the local temperature. The front  face  temperature was  recorded using a 2 colour pyrometer  (Marathon MR1SCSF, Raytek GmbH, Berlin, Germany) and  the  temperature distribution was  recorded using a modified infrared thermal imaging camera (Thermovision A40 FLIR Systems AB, Danderyd, Sweden). The 2 colour pyrometer was capable of recording  temperatures from 1000 to 3000 C and the modified thermal imaging camera could record temperatures up  to 2800 C when combined with  the neutral density filter.2 The aim of the preliminary testing was to rank the UHTC composites according to their oxidation performance and hence the tests were carried out  for 30 s and 60 s. The mass  loss of  the samples after oxyacetylene  torch  testing was recorded and  the depth of erosion was determined from micro-CT  images. The oxidation products were characterised using FEGSEM, EDS, XRD and micro-CT.           3. Results and discussion  The mass of  the Cf preforms before  impregnation and  the mass and bulk density of the composites after impregnation and CVI are summarised in Table 1. The mass increase for the two Hf-based compositions was proportionately  lower  than for  the ZrB2 -based compositions, given  that  the density of  the former is higher. This was due to the larger particle size of the Hf-based powders (<44  \\u242em) compared to ZrB2 (1.5-3  \\u242em), which limited penetration into the carbon fibre preform. Fig. 2a shows a cross-sectional analysis of one of the composites. 15 Cf layers with alternating fibre orientation (0/90) can be distinguished within the composite and the UHTC powder penetrated from one side of the preform to the other. From the higher magnification  images of  the outermost  layers, Fig. 2b and c,  it  Fig. 2.   (a) Powder distributions across the cross section of a UHTC sample and  (b) the top and (c) bottom layers at higher magnifications.  \\x0c', '426   A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432  Fig. 3. 2D micro-CT images of Cf-ZrB2 after impregnation and CVI.  drilled  into  the back  face of  the sample  to  facilitate  the CVI is also clearly distinguishable. The brighter areas  indicate  the presence of UHTC powder;  it  is very evident  that  the UHTC powder has penetrated very well into the preform with the depth 7 mm. of impregnation being  In addition  to  the depth of  impregnation, a good powder distribution is also important to achieve superior high temperature performance. Fig. 4 shows  the EDS mapping on  the cross section of a Cf-ZS20 composite. This  image shows  that good powder mixing was achieved after ball milling, leading to a uniform distribution of  the UHTC powder constituents within  the preform. The time-temperature plots for the UHTC composites tested for 30 s using the oxyacetylene torch are presented in Fig. 5. The composites were  introduced  into  the flame by manually mov2 s. As  ing  the sample stage, a process  that  took  the 2 colour pyrometer can only record  the  temperature above 1000 C,  the moment at which this first reading was recorded on the pyrometer was taken as ‘zero time’ when plotting the time-temperature     Fig. 5. Time-temperature plot   for   the C-C and UHTC composites   tested   for  30 s.           graphs. This, along with the manual monitoring of test duration resulted  in slight differences  in  the  total  test duration  (1-2 s). From the graph it can be seen that the temperature increased at a 500 −1 to  2200 rate of  C s C and then continued to increase slowly  for all  the samples, except  for  the C-C sample where the recorded temperature was much lower than that of all other 1000 −1 and all  samples. The  initial cooling rate was  C s the composites survived the high heating/cooling rates. The  time-temperature plot for  the samples  tested for 60 s  is shown  in Fig. 6. Two C-C samples were  tested at  this  temperature and as  for  the 30 s  test,  the  recorded  temperatures were much lower for the C-C samples even though the same gas flow rates were used for testing all the samples. This is believed to be  Fig. 4. FEGSEM image and EDS mapping on the cross section of a Cf-ZS20 composite showing the powder distribution. Carbon (red), silicon (blue) and zirconium  (green).  \\x0c', 'A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432   427  Fig. 6. Time-temperature plot   for   the C-C and UHTC composites   tested   for  60 s.  due to the absorption of some of the heat by the carbon matrix during ablation.10 Photographic images of the various composites after 30 s oxyacetylene torch testing are shown in Fig. 7. Whilst the effects of oxidation are clear, no appreciable erosion was observed for any of the UHTC-based samples, even though cracking was visible on the front face. An interesting feature observed for the Cf-HfC composite is the lack of adhesion of the oxide layer that formed to the base composite; it fell off during cooling. The main reasons for this is believed to be the build-up of pressure below the oxide  layer due  to CO/CO2 gas  formation and  the absence of formation of any glassy phases during the test. Fig. 8 shows the images of the composites after 60 s of oxyacetylene  torch  testing whilst Fig. 9 shows  the 2D micro-CT images highlighting  the depth of erosion. Table 2 summarises the peak front and back face temperature, mass loss data and erosion depth after  the 30 and 60 s oxyacetylene  torch  tests; note that  the Cf-HfC sample was analysed after  the surface oxide layer had dropped off. The depth of erosion was found to be the  lowest for the Cf-HfB2 system. From Fig. 8, it can be observed that  the C-C sample was eroded over a 20 mm diameter area on  the  front  face of  the sample and  the depth of erosion was 4 mm. The surface erosion of the Cf-ZrB2 and measured to be  Cf-ZS20 composites were similar to each other, with the dam5 mm in diameter and the age mainly focused over an area of  depth of erosion was 4.8 mm and 5.3 mm  respectively.  In  the case of Cf-ZS20-1La composite,  the high  temperature flame penetrated through the dense UHTC powder rich layer attacking the C-C  layer below, resulting  in  increased erosion  to a depth 6.2 mm. The extent of damage was observed  of  to be much lower for  the preforms  impregnated with either HfB2 or HfC. The surface oxide layer formed on the Cf-HfC sample detached within a few minutes of extinguishing  the flame. The presence of molten phases can be seen on the HfB2 and HfC-based composites. On visual inspection, the amount of melting was much lower than that for the ZrB2 -based composites, Fig. 8, because of  the higher melting  temperature of HfO2 (reported  to be as high as 2900 C11 compared to 2715 C for ZrO2 12 ). It is not entirely valid to make a direct comparison between the measured back face  temperatures as  the distances between the front face of the samples, where the flame was focused, and the position where  the  thermocouple was placed were not  the same for all  the samples. Also  the distance between  the flame and the thermocouple changed with time as the high temperature flame eroded the composite. The reactions schemes for the oxidation of the various UHTC constituents used in the composites are given below. 2 O2 →        ZrB2  (s)   + 5   ZrO2 (s)   +   B2O3 (l)   (1)  B2O3 (l)   →   B2O3 (g)   (2)  SiC (s)   + 3  2 O2 (g)   →   SiO2 (l)   +   CO (g)   (passive)   (3)  SiO2 (l)   →   SiO2 (g)   (4)  SiC (s)   +   O2 (g)   →   SiO (g)   +   CO (g)   (active)   (5)  Table 2  Summary of the results after 30 s and 60 s oxyacetylene torch testing.  Composite   Test duration (s)   Peak temperature ( Front face (±150     C)   Weight loss (g)   Erosion depth (mm)     C)   Back face (±10     C)  CC   30   2210   447  491a 477a 582a 548a 530a  0.38   1.0  Cf-ZrB2 Cf-ZS20   30   2560   0.22   Negligible  30   2520   0.18   Negligible  Cf-ZS20-1La   30   2575   0.17  0.77b 0.55b  Negligible  Cf-HfB2 Cf-HfC   30   2625   Negligible  30   2680   Negligible  CC   60   2315   763  857a 723a 877a 918a 847a  1.63   4.0  Cf-ZrB2 Cf-ZS20   60  2590   0.67   4.8  60   2550   0.63   5.3  Cf-ZS20-1La   60   2525   0.74   6.2  Cf-HfB2 Cf-HfC   60   2640   0.57  1.69c  <2.0  60   2530   NM  NM, not measured as the surface layer fell off. a Back face thermocouple placed inside the drilled hole. b Some of the oxide layer fell off during the test.  c  Includes weight of the lost surface layer.  \\x0c', '428   A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432  Fig. 7. Photographs of C-C and Cf-UHTC powder composites after 30 s oxyacetylene torch testing. The diameter of the composites was 30 mm.  Fig. 8. Photographs of C-C and Cf-UHTC powder composites after 60 s oxyacetylene torch test. The diameter of the composites was 30 mm.  \\x0c', 'A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432   429  Fig. 9. Micro-CT images of C-C and Cf-UHTC composites after 60 s oxyacetylene torch test. The images were taken where the depth of erosion was maximum.  +   3B2O3 (l)  + 21  +  + 5  → 1  →  →  LaB6 (s)   4 O2 (g)   2 La2O3 (s)   La2O3 (s)    ZrO2 (s)    La2Zr2O7 (s)  HfB2 (s)   2 O2 (g)    HfO2 (s)    B2O3 (l)  HfC (s)   + 3  2 O2 (s)   →   HfO2 (s)   +   CO (g)   +  (6)  (7)  (8)  (9)  The reaction products formed after the high temperature testing were characterised using XRD as shown in Fig. 10 and the results revealed that the surface layer contained only monocliniczirconia (m-ZrO2 ) for all the ZrB2 containing UHTC composites and monoclinic hafnia (m-HfO2 ) for the HfB2 and HfC containing composites. Of  the various reaction products mentioned  in the  reaction schemes, ZrO2 and HfO2 are stable at >2500 C. B2O3 has a very  low melting point of 450 C and high vapour pressure and hence it will have quickly vaporised at temperatures above 1100 C.13 This will have generated porosity and have accelerated the oxidation process above this temperature. Addition of SiC can significantly improve the oxidation resistance in 1700 the intermediate temperature range from 1200 to  C by forming an SiO2 scale which reduces  the oxygen diffusivity,14 1500 but no SiO2 peaks were observed in the XRD. Above  C, SiC undergoes active oxidation leading to the formation of SiO and CO without  forming a protective  silica  layer and hence                 Fig. 10. XRD analysis of   the oxidation products after 60 s oxyacetylene   torch  testing. All the peaks for the ZrB2 containing compositions correspond to monoclinic zirconia and the all the peaks corresponding to HfB2 and HfC containing compositions correspond to monoclinic hafnia.  it does not offer any additional protection.15 Furthermore,  the formation of water molecules as a result of  the burning of  the acetylene gas can reduce the stability of any SiO2 that does form by generating Si(OH)4 , SiO(OH)2 and/or SiO(OH).16 Addition  Fig. 11. FEGSEM images and EDS pattern on the molten products formed on the surface of Cf-ZS20-1La composite after torch testing.  \\x0c', '430   A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432  Fig. 12. Degradation of carbon fibre after the oxyacetylene torch testing of the C-C composites.     of LaB6 to ZrB2 is  reported  to stabilise  the ZrO2 ,  formed as a result of  the oxidation of ZrB2 ,  in  the  tetragonal phase. The formation of La2Zr2O7 pyrochlore, which has a high melting point (>2300 C), has also been reported.16 From the oxidation testing  it was found  that  the addition of LaB6 actually reduced the high temperature performance of the UHTC composites and XRD analysis did not detect  the presence of any La2Zr2O7 . FEGSEM  images of one of  the molten droplets formed on  the surface of a Cf-ZS20-1La composite showed the formation of  platelet-like  structures  and  an EDS  spectrum  confirmed  the presence of La, Zr,  and O, Fig. 11. So  it  can be  assumed that La2Zr2O7 formed during  the  torch  testing, but melted, segregated and  recrystallised without offering any additional protection. Addition of  lower valence cations such as La3+ to ZrB2 has been reported to increase the oxygen transport through the oxide layer and also to lower the eutectic temperature of the oxide scale leading to accelerated oxidation during high temperature  testing.17 Based on  the  torch  test and XRD results  it can  Fig. 13. Microstructures after 60 s oxyacetylene   testing of a Cf-ZrB2 composite (a) near  droplets which had been molten during the test and (d) high magnification on one of the frozen droplets.  torch   the edge of   the composite, (b) 1-2 mm from flame   tip, (c) frozen  \\x0c', 'A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432   431  Fig. 14. Microstructure after oxyacetylene   torch   temperature and (b) glassy structure formed on the composite      be concluded  that  the addition of SiC or LaB6 did not significantly  improve  the oxidation resistance of UHTC composites when exposed to temperatures >2500 C using the oxyacetylene torch facility. Many interesting microstructures were developed in the composites as a result of  the combination of  the high  temperature, rapid heating/cooling and  thermal gradients during  torch  testing. The carbon fibres of the C-C composites underwent severe degradation, Fig. 12a and b. The surfaces of the fibres were also oxidised leading to pitting, Fig. 12c and d and this type of fibre degradation has been  reported  for Cf composites at elevated temperatures.10,18,19 The microstructures formed after the oxyacetylene torch testing of  the Cf-ZrB2 composites, Fig. 13, revealed  that near  the edges the UHTC particles formed large agglomerates that were not strongly bonded  to one another, Fig. 13a. This  is because the  temperatures  experienced by  these particles were  lower than  in  the flame  tip  region. At 1-2 mm  from  the flame  tip, Fig. 13b,  there was good bonding between  the particles and signs of necking. This area was porous and many cracks are also visible. At  the flame  tip boundary, Fig. 13c,  liquid phases  testing of a Cf-ZS20 composite. (a) Porosity observed on  10 mm from the flame tip.  the frozen droplets which had been molten at   the   test  5  formed  and  large,  \\u242em, grains may be  seen. Fig. 13d  is a higher magnification  image of one of  the droplets  formed after  the  test. This  also  shows  signs of  fusing between  the particles. The microstructures of the Cf-ZS20 and Cf-ZS20-1La composites were more or  less  similar  to  those of  the Cf-ZrB2 composites. Two additional features observed on the surface of the Cf-ZS20 composite are shown  in Fig. 14. The porosity  in what had been molten droplets at higher temperature, Fig. 14a, is believed to be due to the escape of B2O3 , SiO and CO/CO2 gases generated as a result of  the oxidation of ZrB2 and SiC respectively. The formation of glassy microstructures on  the surface, Fig. 14b,  indicated  the formation of borosilicate glasses away from the flame tip, where the sample experienced <2000 C. The microstructure of  the Cf-HfB2 composite subjected  to 60 s oxyacetylene  torch  testing  is shown  in Fig. 15. This composition offered the best high temperature oxidation protection, even  though  the powder  impregnation was not as good as  that achieved  for  the ZrB2-based composites, as discussed earlier. Fig. 15a reveals the presence of what had been molten HfO2 and Fig. 15b reveals fibre degradation  that occurred directly below     Fig. 15. Microstructure after oxyacetylene   torch   temperature and (b) glassy structure formed on the composite   testing of a Cf-ZS20 composite. (a) Porosity observed on  10 mm from the flame tip.  the frozen droplets which had been molten at   the   test  \\x0c', '432   A. Paul et al. / Journal of the European Ceramic Society 33 (2013) 423-432  the flame tip. Fig. 15c shows a high magnification image on one of the frozen droplets and Fig. 15d shows a porous microstruc1-2 mm away from  ture  the flame  tip. It may be possible  to reduce the extent of fibre damage by improving the HfB2 impregnation and experiments are underway  to achieve  this. UHTC powder with finer particle sizes  is required  to  improve further the impregnation.  4. Conclusions  The potential of Cf-UHTC composites  for ultra-high  temperature applications where ablation  is also  relevant has been assessed in this study by preparing composites utilizing a slurry impregnation and carbon CVI route. Based on the high temperature oxidation  testing,  it can be concluded  that  impregnation of Cf preforms with UHTC powders significantly improves the high  temperature oxidation resistance of  the composites compared  to C-C composites. Hf-based UHTC powders offered superior oxidation protection compared  to Zr-based compositions, even though less could be impregnated into the preforms due  to a  larger mean particle size, while  the addition of SiC and LaB6 did not  improve  the oxidation resistance at  the very high temperatures, >2500 C. Of the two Hf-based compounds investigated, HfB2 composites were deemed to have better oxidation performance as  the oxidation products were adherent to  the base composite. The  thermal shock resistance of all  the UHTC composites was found to be excellent.     Acknowledgements  The authors  thank  the UK’s Defence Science and Technology Laboratory (DSTL) for providing the financial support for this work under contract number DSTLX-1000015267 as well as the US Air Force Research Laboratory’s Materials and Manufacturing Directorate for ongoing collaborations.  References     1. Opeka MM, Talmy IG, Zaykoski JA. Oxidation-based materials selection for  2000 C + hypersonic aerosurfaces: theoretical considerations and historical experience. J Mater Sci 2004;39:5887-904.  2. Paul A,   Jayaseelan DD, Venugopal S, Zapata-Solvas E, Binner   J, Vaid hyanathan B, et al. UHTC composites   for hypersonic applications. Am  Ceram Soc Bull 2012;91:22-9.  3. Talmy IG, Zaykoski JA, Opeka MM. Synthesis, processing and properties of TaC-TaB2 -C ceramics. J Eur Ceram Soc 2010;30:2253-63. 4. Han  J, Hu P, Zhang X, Meng S, Han W. Oxidation-resistant ZrB2 -SiC composites at 2200  C. Compos Sci Technol 2008;68:799-806.     5. Sayir A. Carbon fiber reinforced hafnium carbide composite. J Mater Sci 2004;39:5995-6003.  6. Levine SR, Opila EJ, Halbig MC, Kiser JD, Singh M, Salem JA. Evaluation  of ultra-high temperature ceramics for aeropropulsion use. J Eur Ceram Soc 2002;22:2757-67.  7. Levine SR, Opila EJ, Robinson RC, Lorincz JA. Characterization of an ultra high temperature ceramic composite. In: NASA TM-2004-213085. 2004. p.  1-26.  8. Tang S, Deng J, Wang S, Liu W, Yang K. Ablation behaviors of ultra-high temperature ceramic composites. Mater Sci Eng A 2007;465:1-7.  9. Zhao D, Zhang C, Hu H, Zhang Y. 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Active-to-passive   transition   in   the   oxidation  of   silicon carbide and   silicon nitride   in air. J Am Ceram Soc 1990;73:  1540-3.  16. Zhang X, Hu P, Han J, Xu L, Meng S. The addition of   lanthanum hexa boride to zirconium diboride for improved oxidation resistance. Scr Mater 2007;57:1036-9.  17. Hu P, Zhang X, Han J, Luo X, Du S. Effect of various additives on the oxi dation behavior of ZrB2 -based ultra-high-temperature ceramics at 1800  J Am Ceram Soc 2010;93:345-9.     C.  18. Cho D,   Lee   JY, Yoon   BI. Microscopic   observations   of   the   abla tion behaviours of  1993;12:1894-6.  carbon fibre/phenolic   composites.   J Mater Sci Lett  19. Han W, Hu P, Zhang X, Han J, Meng S. High-temperature oxidation at  1900 C of ZrB2 -xSiC ultrahigh-temperature ceramic composites. J Am Ceram Soc 2008;91:3328-34.     \\x0c']"
},{
  "_id": 275,
  "PDF": "Ultra-high temperature ceramics Materials for extreme environments.pdf",
  "Text": "['Scripta Materialia 129 (2017) 94-99  Contents lists available at ScienceDirect  Scripta Materialia  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / s c r i p t a m a t  Viewpoint article  Ultra-high temperature ceramics: Materials for extreme environments  William G. Fahrenholtz ⁎, Greg E. Hilmas  Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65049, United States  a r t i c l e  i n f o  a b s t r a c t  This paper identifies gaps in the present state of knowledge and describes emerging research directions for ultrahigh temperature ceramics. Borides, carbides, and nitrides of early transition metals such as Zr, Hf, Nb, and Ta have the highest melting points of any known compounds, making them suitable for use in extreme environments. Studies of synthesis, processing, densification, thermal properties, mechanical behavior, and oxidation of ultra-high temperature ceramics have generated a substantial base of knowledge, but left unanswered questions. Emerging research directions include testing/characterization in extreme environments, composites, computational studies, and new materials. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  Article history:  Received 1 September 2016 Received in revised form 12 October 2016 Accepted 13 October 2016 Available online 20 October 2016  Keywords:  Ceramics Extreme environments Borides Carbides Nitrides  1. Introduction  Ultra-high temperature ceramics (UHTCs) are an emerging class of materials that have the potential for use in extreme environments [1,2]. Most often, UHTCs are defined as compounds that have melting points above 3000 °C (Fig. 1) [3]. Several alternative definitions have been proposed with the most pragmatic being that UHTCs are ceramic materials that can be used for extended times at temperatures above 1650 °C [4]. However, none of these definitions captures the wide range of extreme conditions in which UHTCs may be used. Chemically, nearly all UHTCs are binary compounds in which boron, carbon, or nitrogen combine with one of the early transition metals (TMs) such as Zr, Hf, Ti, Nb, and Ta [5]. The strong covalent bonds between the TMs and B, C, or N result in compounds with high hardness, stiffness, and melting temperature [6]. These compounds also exhibit higher electrical and thermal conductivities than oxide ceramics due to varying degrees of metallic bond character. This intriguing combination of metal-like and ceramic-like properties allow UHTCs to survive extreme temperatures, heat fluxes, radiation levels, mechanical loads, chemical reactivities, and other conditions that are beyond the capabilities of existing structural materials. Reports of the synthesis of UHTCs date back to the late 1800s [7,8], but technological interest in these materials grew dramatically during the space race that started in the late 1950s and extended through the 1960s. At that time, groups in the Soviet Union and the U.S. were searching for materials to be used in rocket motors, heat shields, and structural components for the first generation of spacecraft [9,10].  ⁎  Corresponding author. E-mail address: billf@mst.edu (W.G. Fahrenholtz).  http://dx.doi.org/10.1016/j.scriptamat.2016.10.018 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.  These vehicles were going to encounter heat fluxes and temperatures for durations that were beyond the capabilities of traditional approaches such as transpiration cooling, heat sinks, and convective/radiative cooling [11]. The technological challenges motivated a tremendous body of research conducted by scientists including G.V. Samsonov, H. Nowotny, E.V. Clougherty, and L. Kaufman. Publications produced by these researchers and their collaborators document fundamental understanding of the bonding, electronic structure, thermodynamic properties, mechanical behavior, phase equilibria, and oxidation response of UHTCs. After this initial surge of interest, research on these materials waned for several decades. Interest in UHTCs was renewed in the late 1980s, again due to aerospace applications such as scramjet propulsion, hypersonic aerospace vehicles, and advanced rocket motors [4,12-15]. For example, in the U.S., the National Aeronautics and Space Administration (NASA) and the U.S. Air Force built and flew vehicles as a means to test technologies needed for sustained hypersonic flight (Fig. 2). One of the technological challenges was presented by the designs, which included sharp leading edges and control surfaces. At hypersonic speeds, the bow shock at these sharp features produces temperatures in excess of 2000 °C, heat fluxes of hundreds of Watts per square cm, and highly reactive dissociated gas species [16]. Diboride-based UHTCs were considered for these applications based on the combination of thermal conductivity, strength at elevated temperature, and oxidation resistance [17], while carbides and nitrides have received more attention for components such as nozzle throats, divert/attitude control thrusters, and nozzle liners where higher thermal and mechanical loads are encountered [18]. In this paper, continuing challenges in areas of recent research activity are identified followed by an assessment of emerging trends and an outlook for future research.  \\x0c', 'W.G. Fahrenholtz, G.E. Hilmas / Scripta Materialia 129 (2017) 94-99  95  composites. Most of the liquid/polymer processes reported to date have involved two steps whereby the precursors decompose to an oxide with or without carbon/boron, which is then reacted to form the desired UHTC phase (Fig. 3) [22]. Some recent advances have surmounted this obstacle by directly synthesizing borides or carbides in a single step [23]. Continuing research challenges in synthesis include synthesizing higher purity UHTC powders, identifying mechanisms and kinetics of synthesis reactions, and designing liquid/polymer precursors that have higher ceramic yields and are less sensitive to degradation under ambient conditions. Most often, UHTCs are processed as dry powders using conventional forming methods such as dry pressing. A limited number of studies have reported novel aspects of dispersion [24], colloidal forming [25], and plastic forming methods [26,27], revealing fundamental aspects of the surface chemistry of UHTCs while also enabling near-net shape forming. Processing and synthesis have substantial synergy whereby improvements in synthesis that control the size, shape, or surface chemistry of UHTC powders impact the ability to form stable dispersions and develop desired rheological properties in suspensions, pastes, or plastic masses. Very few studies have addressed the fundamental processing science of UHTCs resulting in opportunities to explore control of powder surface chemistry, produce slurries and pastes with high solids loadings, develop additive manufacturing processes, and devise aqueous processing routes that do not increase the oxygen impurity content of the resulting ceramics. Sintering of non-oxide ceramics, and particularly UHTCs, involves a competition between densifying and coarsening mechanisms [28]. The presence of oxygen and other impurities promotes coarsening mechanisms that produce porosity trapped within grains, exaggerated grain growth, or other phenomena, which prevent ceramics from reaching full relative density (Fig. 4) [29]. Finer starting particle sizes can enhance sintering, but oxygen contents must be below certain critical levels for full densification. While pressure-assisted methods such as hot pressing, hot isostatic pressing, and spark plasma sintering are the most common methods for densification, pressureless sintering can be accomplished in some systems through the use of additives [30-32]. Despite practical knowledge for densification, only a few publications have reported detailed information on the fundamental transport mechanisms for densification or the kinetics of sintering processes leaving ample room for additional basic research in these areas. A multitude of studies have described the room temperature thermal and mechanical properties of UHTCs with the most important properties being strength, hardness, elastic constants, thermal conductivity, and fracture toughness. Most of these studies have reported properties for a single ceramic without identifying fundamental factors related to composition, microstructure, porosity, impurities, etc. that control observed behavior. Presentation of properties without understanding of structure-property relations presents a particular challenge for UHTCs. A single report of poor properties or sub-standard performance in a ground test can turn potential users against an entire class of materials, even if the material in question failed due to extrinsic factors that could have been addressed had the root cause of failure been studied. Systematic studies have elucidated microstructure-mechanical property relationships and identified strength-limiting flaws to the extent that ceramics with desired properties can be designed by controlling grain size for nominally pure, fully dense ZrB2 or SiC cluster size in ZrB2-SiC ceramics. Studies have also begun to address the fundamental factors that are responsible for the wide variety of room temperature thermal conductivities reported for ZrB2-based ceramics (Fig. 5) and other UHTCs [33-36]. Additional research is needed to isolate fundamental factors that control thermal and mechanical behavior for compositions other than ZrB2 or ZrB2-SiC, examine elevated temperature structureproperty relationships, and report intrinsic properties of UHTCs. Oxidation behavior, or, more broadly, resistance to aggressive chemical environments, is critical in the extreme environments that might be encountered by UHTCs include oxidation above 1600 °C, reaction  Fig. 1. Melting temperature for elements and compounds from different families of materials showing that very few materials have melting temperatures above 3000 °C. Reprinted by permission from Fahrenholtz et al. [3].  2. Continuing challenges  Research on UHTCs during the past 25 years has furthered the state of knowledge in several important areas including synthesis, processing, densification, thermomechanical properties, and oxidation. Despite these advances, key gaps in the understanding of UHTCs remain. Research groups in the U.S., China, Japan, the U.K., Italy, India, and other countries have on-going efforts focused in these areas. Most commercial UHTCs are synthesized by carbothermal reduction of the corresponding oxide [19,20]. Limitations of this method include the relatively coarse particle sizes and substantial levels (i.e., 0.1 wt% or higher) of impurities including other metals, oxygen, nitrogen, and/ or carbon. Despite these limitations, a majority of the recent studies of UHTCs have utilized commercial powders. Alternative synthesis methods can produce powders with finer particle sizes, higher purities, or controlled shapes [21]. Processes based on solid-state reactions have been the most popular routes, but processes utilizing liquid or polymer precursors have been pursued to enable deposition of coatings or polymer infiltration and pyrolysis (PIP) processing of UHTC-matrix  Fig. 2. Photograph of the X-51 flight vehicle mounted on the wing of a B-52 prior to a test flight. The X-51 holds the record for the longest duration hypersonic flight by an aircraft. Note the heads of engineers on the lower left of the image for scale. Image is public domain from NASA.  \\x0c', '96  W.G. Fahrenholtz, G.E. Hilmas / Scripta Materialia 129 (2017) 94-99  Fig. 3. Schematic of the two step process for the formation of ZrB2-SiC ceramics from a polymeric precursor showing that ZrB2 is formed by decomposing the polymeric precursor to ZrO2 and C followed by reaction with B2O3 to form ZrB2. Reprinted by permission from Cao et al. [22].  dissociated air behind a bow shock, or corrosion in the exhaust plume of a solid rocket motor. Reactions between TM compounds of B, C, and N and oxygen are highly favorable at all temperatures. Additives commonly used to improve oxidation resistance include the addition of Sirich compounds with SiC or MoSi2 being the most common. These additives improve oxidation resistance by promoting passive oxidation through the formation of a silica-rich scale. However, silica scales are only effective in certain regimes of temperature, total pressure, and external gas velocity. Additives such as SiC that improve oxidation resistance in conventional oxidation tests can have a detrimental effect in  extreme environments by generating gaseous reaction products through active oxidation [37]. Under the most extreme conditions, the TM oxide provides the protection. In general, borides have superior oxidation resistance at temperature up to ~ 1700 °C while carbides perform better at higher temperatures. Significant progress has been made toward understanding both the thermodynamics and kinetics of the oxidation of diboride-based UHTCs (Fig. 6) [38]. In addition, oxidation behavior of the carbides has also been reported. The main challenges for oxidation studies are to understand behavior in extreme environments that are relevant to different applications, predicting the effects of composition and microstructure on oxidation without prior experimental studies, and designing UHTC structures that are thermomechanically stable and resistant to oxidation.  3. Emerging trends  Several exciting new research directions are emerging for UHTCs. These studies draw on the base of knowledge established in historic studies or recent research to address challenges that stand in the way of UHTCs reaching their full potential for use in extreme environments. For this paper, emerging research has been grouped into theme areas of testing/characterization in extreme environments, composites, computational studies, and new materials. While these themes do not capture all of the efforts that are in progress, they provide a snapshot of some of the significant research challenges for UHTCs.  Fig. 4. Comparison of microstructures of ZrB2 + 30 vol% SiC ceramics after pressureless sintering showing (left) low relative density (~ 60%) and large particles (N 50 μm) with entrapped porosity produced by sintering without sufficient additives and (right) fully dense ceramic sintered with the addition of B4C and SiC. Note the difference in magnification for the images. Reprinted by permission from Zhang et al. [29].  Fig. 5. Reported values for thermal conductivity as a function of temperature for a series of nominally pure ZrB2 ceramics showing that room temperature thermal conductivity for ZrB2 can vary from ~ 30 W/m·K to N 130 W/m·K. Compiled using data from Guo et al. [33], Zhang et al. [34], Harrington et al. [35] and McClane et al. [36].  \\x0c', 'W.G. Fahrenholtz, G.E. Hilmas / Scripta Materialia 129 (2017) 94-99  97  Fig. 6. Schematic of the oxidation of a TMB2 ceramic containing SiC showing the formation of a scale containing the TM oxide and a borosilicate glassy layer, which recedes during oxidation at elevated temperatures. Reprinted by permission from Parasarathy et al. [38].  A vast majority of the characterization and testing of UHTCs is carried out under ambient conditions (i.e., near room temperature and atmospheric pressure) even though applications of these materials will expose them to extreme environments such as temperatures of 2000 °C or above, heat fluxes of hundreds of W/m·K, dissociated gases, etc. Unfortunately for the research community, much of the testing that has been done in relevant environments is restricted or classified since those environments and the test results could reveal the capabilities and weaknesses of defense systems or information that could be used to proliferate nuclear weapons technology. The promise of improved performance under extreme conditions is often made by extrapolating data collected at near-ambient conditions without accounting for potential non-linear behavioral trends, which could be caused by phenomena such as phase changes, creep, softening of grain boundary/impurity phases, sub-critical crack growth, or stress-induced microcracking. A few research laboratories around the world have established the ability to test materials under extreme conditions. For example, our laboratory at the Missouri University of Science and Technology can test mechanical properties (strength, elastic modulus, fracture toughness), thermal properties (thermal diffusivity and heat capacity), and electrical conductivity from room temperature up to 2000 °C or higher [39-44]. Likewise, research groups at the Harbin Institute of Technology [45,46], Imperial College London [47], University of Birmingham [48], and University of Arizona [49] have all established testing capabilities that reproduce one more types of extreme environments. In addition, a team at University of California, Berkeley has developed a unique chamber that allows in-situ observation of mechanical tests at ultra-high temperatures [50]. Without a doubt, the ability to produce and sustain extreme conditions in test facilities is difficult, particularly when attempting to create combined environments to evaluate the influence of multiple stresses (e.g., mechanical load and heat flux). Fundamentally, most sensors do not operate in extreme conditions so accurate characterization of environments and material responses is difficult. The ability to test and characterize materials under extreme conditions is needed not only to understand the fundamental behavior under these conditions, but also to enable collection of data to verify simulations of performance for UHTCs in specific applications that involve one or more extreme environments. Several laboratories around the world have the ability to simulate conditions associated with hypersonic flight and atmospheric re-entry using plasma wind tunnels, inductively coupled plasma systems, arc heaters, or similar systems [51-59]. While some of these are open to academic researchers, results tend to be restricted if they are generated at facilities operated by government agencies and companies that reproduce conditions for specific flight profiles. To overcome issues of cost and access, groups have built oxy-acetylene torch rigs for testing under high heat flux/ablation conditions [60-63]. Both types of tests are vital to demonstrate that UHTCs can survive the conditions associated with hypersonic flight, rocket propulsion, atmospheric re-entry, etc.  However, the reported studies tend to evaluate a single material under a specific set of conditions, which does not produce significant insight into the fundamental understanding of the material response to the extreme conditions. For example, ablation or recession rates are often calculated for a UHTC based on a single exposure condition, but the response of UHTCs to these environments is complex and inherently non-linear. Hence, calculation of an ablation rate from a single exposure is extremely suspect. Additional limitations of these tests are that they do not directly evaluate any of the fundamental properties of the materials and the environments, though extreme, do not reproduce those that will be encountered in use. As a result, tests in relevant environments are under-utilized by basic researchers, creating an opportunity for fundamental research on the complex responses (i.e., coupled thermal, mechanical, chemical, etc.) exhibited by materials in extreme environments. Continuous fiber reinforced UHTCs are needed because UHTCs are brittle and exhibit ceramic like mechanical behavior. As a result, exposure to extreme temperature, heat fluxes, strain rates, etc., can produce catastrophic failure. For example, monolithic ceramics tested in simulated atmospheric re-entry conditions often fail during testing due to thermal shock [59]. Composites in which UHTC matrices are reinforced by continuous fibers are attractive because of the improved fracture toughness and non-catastrophic failure behavior. Several groups have begun to investigate UHTC-based composites [64-70]. Some of the continuing technical challenges include formulating liquid/polymer-based precursors for UHTC matrices, producing fibers capable of withstanding extreme conditions, synthesizing high purity UHTC powders with controlled particle sizes for infiltration processing, and increasing the elastic modulus of available fibers to promote load transfer from the brittle UHTC matrix to the reinforcing fiber. Computational materials research methods have been applied to some UHTC systems. Several authors have predicted fundamental properties, thermodynamic stability, stable crystal structures, and deformation mechanisms for different boride, carbide, and nitride systems. For example, hardness and elastic constants have been predicted computationally in the search for ultra-hard, incompressible materials [6,71-74]. Some initial studies of phase stability and deformation have been reported along with simulations related to thermal properties [75-78]. Challenges for computational investigations include defining accurate interatomic potentials for UHTC systems, identifying properties that reflect intrinsic behavior of the materials rather than extrinsic factors for model validation, and extending simulations to ultra-high temperatures. More robust simulations of fracture behavior and transport properties are also needed to understand mechanisms at length scales ranging from atoms to the mesoscale. The family of UHTCs being studied at the present time are essentially the same materials that were examined during the space race of the 1950s and 60s. Discovery of new materials is needed to expand the number of UHTCs and thereby broaden the number of potential  \\x0c', \"98  W.G. Fahrenholtz, G.E. Hilmas / Scripta Materialia 129 (2017) 94-99  applications for UHTCs. The likelihood that a large number of new UHTC compounds remains undiscovered seems extremely low, which means that research on new UHTCs will more likely follow the development of advanced structural metals [79,80]. For example, opportunities exist to design second phase additions that could enable transformation toughening or inhibit creep deformation at elevated temperatures while simultaneously improving oxidation resistance. Likewise, additives that dissolve into the UHTC matrix and form solid solutions are another method by which properties might be manipulated. Our group has demonstrated that oxidation behavior at intermediate temperatures can be improved by adding transition metals such as W, Ta, or Nb, but solution additives may also benefit strength and/or fracture toughness at room or elevated temperatures [81,82]. Rather than conventional experimental approaches, the search for new UHTCs is likely to be driven by computational means and will likely require novel approaches such as exploring entropy-stabilization [83] or other methods that have not yet been espoused. Such approaches have the potential to revolutionize the search for, not just new UHTCs, but all types of new materials. Several new applications will also motivate future research on UHTCs. Whereas most of the discussion above described aerospace applications, UHTCs are also candidates for use in applications such as advanced nuclear fission reactors [84], high temperature electrodes for metal refining [85], high power-density microelectronics [86], concentrated solar power [87], fusion energy systems [88], and many others. In particular, nuclear applications have gained significant attention due to needs for accident-tolerant fuels and cladding, non-oxide fuel pellets, inert matrix fuels, waste separation, and moderators [89]. A growing number of possible applications should lead to additional research in this area. Some of the anticipated challenges are cost, minimizing impurities, producing dense materials, and estimating performance lifetime.  4. Summary and outlook  Significant research effort has gone into studying the fundamental properties of UHTCs. Revealing the remarkable properties of UHTCs continues to broaden the possible range of potential applications. The initial motivation for UHTC research was centered on aerospace applications due to needs for materials with higher melting temperatures, improved thermal conductivity, mechanical stability, and resistance to the extreme heat fluxes and chemically reactive plasmas encountered during hypersonic flight and atmospheric re-entry. Presently, the range of potential applications for these materials is broadening to include the fields of energy (nuclear fission and fusion, energy harvesting, concentrated solar power), materials processing (high temperature electrodes, high speed machining tools, molten metal containment), microelectronics (conductors, barrier layers, lattice matched substrates), and others. The lure of revolutionary advances in performance and the need to withstand increasingly extreme operating conditions will continue to motivate research on UHTCs. Research efforts should evolve from the present focus on synthesis, processing, densification, structure-property relationships and oxidation to tackle challenges associated with performance in extreme environments, fiber-reinforced composites, computational studies, and new materials. By about 2020, the second wave of UHTC research will have been active for nearly 30 years, which is significant because this is the typical time required for a new class of materials to find its initial commercial applications. As UHTCs reach this milestone, potential funding sources will likely expand from the predominance of government funded efforts focused on basic research to include increasingly applied research and development efforts funded by companies and industrial consortia. Hence, the global research community that studies UHTCs should continue to grow as UHTCs transition from laboratory studies to real-world applications.  Acknowledgement  This work was supported by the Ultra-High Temperature Materials Program in the Office of Naval Research through grant N00014-16-12303.The guidance of program manager Dr. Eric Wuchina is acknowledged. 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  "_id": 276,
  "PDF": "Ultra-high temperature oxidation of a hafnium carbide-based solid solution ceramic composite.pdf",
  "Text": "['Corrosion Science 80 (2014) 402-407  Contents lists available at ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  Ultra-high temperature oxidation of a hafnium carbide-based solid solution ceramic composite  David W. Lipke a,⇑ , Sergey V. Ushakov b, Alexandra Navrotsky b, Wesley P. Hoffman a  a Air Force Research Laboratory, AFRL/RQRC, Edwards AFB, CA 93524, USA b Peter A. Rock Thermochemistry Laboratory and NEAT ORU, University of California at Davis, Davis, CA 95616, USA  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 30 May 2013 Accepted 29 November 2013 Available online 4 December 2013  Keywords:  A. Ceramic C. High temperature corrosion C. Oxidation  1. Introduction  Hafnium carbide-based solid solution ceramics with yttria additions were consolidated via hot pressing utilizing chromium carbide as a transient liquid sintering agent. Selected samples were oxidized at temperatures exceeding 2200 °C under high velocity airflow to determine their potential for service in ultrahigh temperature environments. Cross-sectional examination revealed a complex multi-layer scale comprising a highly porous exterior region, a dense multi-phase interior region, and an oxygenand carboncontaining interlayer consisting of nanocrystalline graphite in an amorphous or nanocrystalline oxide matrix. Outward diffusion of mobile cations is speculated to play an important role in establishing the observed scale morphology.  Ó 2013 Elsevier Ltd. All rights reserved.  The discovery and development of ultra-high temperature (i.e., greater than 2000 °C) oxidation-resistant materials is expected to enable many future technologies such as protective coatings for hypersonic and propulsion-related aerosurfaces [1,2]. Few known materials possess the requisite refractoriness to withstand such temperatures, let alone the simultaneous corrosive action of oxidizing species. Investigators in this field have expended great effort to form an understanding of the effects of chemical composition and processing-microstructural relationships on the oxidation behaviors of two classes of relatively oxygen impermeable materials: silica/silicate-formers and iridium/rhenium-based compounds [3]. Due to the onset of active oxidation (i.e., wherein the primary oxidation products are gaseous rather than condensed phases) silica/silicate-formers are unsuitable for extended exposures to ultra-high temperatures [4-7], while the cost and scarcity of iridium/rhenium-based compounds makes their widespread use impractical. The search for viable alternatives has led to an interest in hafnium-bearing materials (e.g., HfC, HfB2, Hf-based alloys, etc.) whose primary oxidation product hafnia is amongst the most refractory oxides [8-11]. Under certain conditions, the oxidation of hafnium carbide-based materials has been observed to produce an oxygenand carbon-containing interlayer between the unoxidized carbide and oxide scale [2,12-16]. Oxidation kinetics studies  ⇑ Corresponding author. Present address: Kazuo Inamori School of Engineering at Alfred University, Alfred, NY 14802, USA. Tel.: +1 (607)871 2729.  E-mail address:  lipke@alfred.edu (D.W. Lipke).  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.11.049  seem to indicate that this oxide-carbide interlayer constitutes a diffusion barrier whose diffusion coefficient is at least an order of magnitude lower than that of hafnia under similar conditions of temperature and oxygen partial pressure [15]. Thermodynamic arguments impose strict limits on the equilibrium oxygen partial pressure at the oxide-carbide interface depending on the local carbon activity. For example, the fourphase mixture embodied by the net reaction HfO2 + C = O2(g) + HfC would only be in equilibrium at an oxygen partial pressure of 4.6 \\x02 10\\x0011 atm at 2200 °C, assuming all solid phases exist in their pure, crystalline states at unit activity [17]. Thus, an effective secondary oxygen barrier (e.g., in the form of a protective overlayer) may also be requisite for the chemical stability of the interlayer. Notwithstanding issues surrounding thermochemical instability of silica/silicate-formers, potential lessons from the oxidation scale morphology of state-of-the-art HfB2/SiC ceramic composites might be applied to the design of oxidation-resistant hafnium carbide-based ceramics. For example, the oxidation scale formed upon HfB2/SiC consists of a porous hafnia framework that lends mechanical integrity to the scale in high shear environments, infilled by a hafniumand boron-containing silicate liquid that acts as the basis for protection against oxygen ingress, sealing cracks and pores while allowing the escape of back-pressured product gases [18- 21]. This self-healing aspect is a highly sought after feature for protective coatings enabled by the inclusion of a flowable but not overly inviscid phase [22,23]. In this preliminary investigation, we sought to develop an ultrarefractory analogue to the ‘porous framework/liquid infill’ morphology of HfB2/SiC oxidation scales via the introduction of several additional components to a hafnium carbide-based solid solution  \\x0c', 'D.W. Lipke et al. / Corrosion Science 80 (2014) 402-407  403  ceramic whose oxidation products might serve to induce limited liquid formation. Chromium carbide (Cr3C2) was identified as a promising transient liquid phase sintering agent to facilitate the formation of homogeneous solid solutions [24,25]. The peritectic decomposition of pure Cr3C2 into a carbon-saturated chromium liquid and free carbon is known to occur at 1826 °C and, more importantly, available ternary C-Cr-Hf, C-Cr-Ti, and C-Cr-Ta phase diagrams indicate the existence of eutectic compositions with substantial carbide solubilities along the pseudo-binary joins of Cr-HfC, Cr-TiC, and Cr-TaC [26]. Since these constituent metal oxidation products possess far lower melting points than their unoxidized carbide counterparts, anticipated solidus depression resulting from chromium substitution into the hafnium-carbide based solid solution ceramic composite was not considered to be overly disadvantageous. Here we report on the synthesis and ultra-high temperature oxidation behavior of a hafnium carbidebased solid solution ceramic composite consolidated via hot pressing utilizing Cr3C2 as a transient liquid phase sintering agent.  2. Materials and methods  Hafnium carbide (Pure Tech Inc., 99.5%), tantalum carbide (Pure Tech Inc., 99.5%), a 60 mol.% titanium carbide/40 mol.% chromium carbide mixture (Inframat Advanced Materials, 99.9%) and yttria (Alfa Aesar, 99.9%) powders were utilized asreceived without further processing. Manufacturer-stated median particle sizes were less than 5 lm which were verified via scanning electron microscopy (SEM). Batch compositions were dry blended (Turbula T2F shaker-mixer) and loaded within a 25.4 mm diameter graphite die designed for a vacuum hot press (Thermal Technologies Inc. Astro model HP20 series). The charge was heated to 1500 °C at which point a pressing load of 30 MPa was applied. This load was maintained upon heating to a maxi1800 °C and continued to mum processing temperature of be held until the recorded ram travel was less than 0.001 mm within any 60 s interval. Hot pressed specimens were characterized via X-ray diffraction (XRD) (Scintag Inc. XGEN-4000 with Cu anode), SEM (FEI Quanta 600), energy dispersive spectroscopy (EDS) (IXRF Systems model 550i), Vickers microhardness (Shimadzu HMV), and subjected to Archimedes density determination via water displacement method. Two compositions were selected for ultra-high temperature testing at the Laser Hardened Materials Evaluation Laboratory (LHMEL) (Wright-Patterson Air Force Base, Dayton, OH). Specimens (approximately 10 mm thick, 25 mm diameter) were mounted in a split two-piece graphite holder under compressive spring loading using four 5 mm diameter zirconia beads as contact points and subjected to laser heating (IPG 10 kW 1.07 lm fiber optic laser) for up to 30 min. The specimen holder was mounted within a subsonic wind tunnel that imposed a 90 m/s air flow across the illuminated surface to simulate the high shear conditions encountered by aerosurfaces. Oxidized specimens were cross-sectioned, polished, and further analyzed via SEM, Raman spectroscopy (Renishaw RM1000 using 515 nm argon laser excitation), and electron backscatter diffraction (EBSD) (conducted by EBSD Analytical Inc., Lehi, UT). The oxidation scale was also destructively sampled via diamond abrasion with recovered material characterized via XRD (Inel Equinox with Co anode). Finally, compositional data was obtained via wavelength dispersive spectroscopy (WDS) (Cameca SX-100); however, significant quantification error may have been introduced due to the use of available standards that did not closely match specimen composition. As a result, reported WDS-derived values have been truncated to the ones digit and are considered ‘semi-quantitatve’.  3. Results and discussion  Preliminary consolidation trials empirically revealed that batch compositions comprising nominally 20 vol.% Cr3C2 could allow for substantially complete powder consolidation within 30 min at 1800 °C under 30 MPa pressure. XRD analyses, as represented in Fig. 1, indicated that complete carbide solid solutions were achieved, presumably hastened by the dissolution-reprecipitation action of the transient liquid phase sintering agent. Compositional analyses and selected property data are presented in Table 1 for each test specimen. Six specimens (three each of two batch compositions) were thus tested at the LHMEL facility under various laser power configurations. Due to the limited number of trials, only three specimens (A1, B1, and B2) were successfully tested under controlled conditions. Each of these three specimens exhibited severe lateral cracks during initial heatup, which we speculate to be due to a combination of thermal shock and compressive failure owing to insufficient compliance of the sample holder against thermal expansion of the specimen. Specimen A1 experienced a maximum front surface temperature of 2190 °C and exhibited catastrophic surface recession due to excessive liquid formation, presumably associated with the greater tantalum carbide batch content whose primary oxidation product Ta2O5 melts at 1872 °C and is not known to form any high melting point tantala-rich compounds with hafnia, titania, chromia, or yttria. As a result, the ‘A’ (Ta-rich) batch composition was deemed unsuitable for ultra-high temperature applications. Specimen B1 demonstrated significant surface recession upon experiencing a peak surface temperature of 2350 °C, seemingly exceeding the maximum temperature capability of the ‘B’ batch composition. Specimen B2 reached an average front surface temperature of 2210 °C with heating profiles for the front and back surfaces shown in Fig. 2 and was observed to form a compact oxidation scale whose structure and composition was analyzed in detail and discussed below. Cross-sectional examination of specimen B2 revealed an adherent, complex multi-layer oxidation scale consisting of a highly porous outer region and relatively compact interior, including an oxygenand carbon-containing interlayer between the unoxidized carbide composite and fully oxidized region. Several thin transverse cracks were observed to span the compact interior region (with endpoints indicated by arrows in Fig. 3a). The absence of enhanced oxidative corrosion in the vicinity of the cracks supports the notion that they likely formed upon cooldown. If the oxidation scale and unoxidized specimen together act as would a bilayer structure, and assuming the average linear coefficient of thermal  Fig. 1. XRD pattern of specimen B2 post-hot press consolidation at 1800 °C for 40 min. under 30 MPa applied load (y = Y2O3, ss = solid solution carbide).  \\x0c', '404  D.W. Lipke et al. / Corrosion Science 80 (2014) 402-407  Table 1 Compositional analyses and selected property data for test specimens.  ID  A1 A2 A3 B1 B2 B3  Y2O3 (mol.%)a  Hafnium carbide-based solid solution  HV1 (GPa)d  q (g/cm3)e  5.2 ± 0.8 5.5 ± 0.5 5.2 ± 0.6 5.1 ± 0.5 5.3 ± 0.5 5.2 ± 0.8  Compositionb  (Hf0.68Ti0.17Ta0.12Cr0.02)C (Hf0.71Ti0.15Ta0.12Cr0.02)C (Hf0.67Ti0.14Ta0.12Cr0.07)C (Hf0.79Ti0.13Ta0.05Cr0.04)C (Hf0.79Ti0.12Ta0.04Cr0.05)C (Hf0.77Ti0.13Ta0.05Cr0.05)C  Lattice parameter (Å)c  4.541 ± 0.002 4.551 ± 0.002 4.5477 ± 0.0007 4.5723 ± 0.0004 4.5750 ± 0.0007 4.5694 ± 0.0005  18.4 ± 0.9 18.8 ± 0.8 18.5 ± 0.6 18.5 ± 0.7 18.2 ± 0.7 18.0 ± 0.5  10.29 ± 0.01 10.43 ± 0.01 10.40 ± 0.01 10.42 ± 0.01 10.30 ± 0.01 10.31 ± 0.01  a Molar content relative to overall composite with the balance comprising the hafnium carbide-based solid solution. b Based on semi-quantitative WDS analyses. c Determined by XRD using Y2O3 as an internal standard. d Vickers microhardness using 1 kg load. e Archimedes density using water displacement method.  scale thickness of 1360 lm in the centre of the laser spot. The apparent porosity varied from greater than 60% in the outer scale layers to less than 10% in the interior region. EBSD analyses (shown in Fig. 3b) within the outermost porous layer (4) could be indexed to cubic symmetry, while semi-quantitative WDS analyses established the composition as (Hf0.85Y0.10Ta0.04Ti0.01)O2. The next porous layer (3) has been identified as exhibiting monoclinic symmetry with composition (Hf0.81Ti0.09Ta0.05Y0.04Cr0.01)O2. XRD analyses (shown in Fig. 3c) confirmed the presence of cubic and monoclinic phases in the porous outer region, though apparently part of layer (4) had mixed with layer (3) during abrasive sampling for XRD. Beneath this porous outer region exists a relatively compact multi-phase layer displaying gradients in composition and phase distribution. The exterior portion of this layer, denoted (2)ext., is believed to consist of orthorhombic HfTiO4 and cubic stabilized hafnia, as supported by EBSD analysis and EDS/WDS compositional analyses. In this region the grain size was sufficiently coarse so as to be resolved via WDS elemental mapping (Fig. 3d), which revealed that HfTiO4 was depleted in tantalum, while cubic stabilized hafnia was enriched in yttrium and chromium content. Gradients in chromium and titanium concentration across layer (2)ext. could be detected via WDS analyses. As shown in Fig. 3e, titanium content sharply increases then plateaus in association with HfTiO4 precipitation, whereas chromium content continually increases. The chromium content hereafter sharply drops in the across the layer (2)ext.-(3) interface, presumably due to enhanced chromia volatility and gas phase transport through the open pore network characteristic of the non-protective layers (3) and (4). Moving deeper through this layer, the average phase size apparently decreases until it can no longer be resolved via WDS mapping, yet the larger crystallites could be detected via EBSD. In this interior region, denoted (2)int., an increased fraction of monoclinic hafnia was observed via XRD, though since it could not be detected via EBSD the crystallite size may be below \\x18100 nm (the approximate instrument detection limit). Nanocrystallinity of monoclinic hafnia in this region is further supported by Scherrer the [\\x001 1 1] and [1 1 1] reflecanalysis of the peak broadening of tions, from which a crystallite size of 20 nm could be calculated after accounting for instrument broadening. The final layer (1) beneath this multiphase region consists of an oxygenand carbon-containing interlayer, as shown in Fig. 4. High resolution EBSD analyses of this region (Fig. 4a) identified 1-5 lm particles of undissolved yttria and few 100-400 nm monoclinic crystallites; however, the preponderance of the layer could not be assigned any crystalline status. WDS analyses (Fig. 4b) revealed gradients in oxygen and carbon content across the oxide-carbide interlayer and an average empirical formula of Cr0.05)O1.7C, excluding the presence of discrete undissolved yttria particles. Raman spectroscopy (Fig. 4c) of the interlayer displayed  (Hf0.7Ti0.2Ta0.07 Fig. 2. Front and back surface temperatures of specimen B2 during ultra-high temperature oxidation testing as averaged from multiple pyrometers.  expansion of the oxide scale is greater than that of the carbidebased specimen, then the thermal mismatch stress upon cooldown would be tensile at the exterior surface and compressive within the carbide specimen. Thus, the crack likely originated at the exterior surface and terminated at the oxide-carbide interlayer. Whether or not the interlayer is inherently crack tolerant remains unknown, though such a capability would be highly beneficial for thermal cycling applications. Compositional and crystallographic analyses (shown in Fig. 3) were able to distinguish at least four distinct product layers of varying porosity and phase content; however, we would like to preface discussion of characterization results by highlighting some of the challenges of phase discrimination in this system. Pure hafnia exists as a monoclinic phase at room temperature, tetragonal phase above 1820 °C, and cubic phase above 2627 °C. Substitution of aliovalent or large tetravalent cations has been shown to stabilize high temperature tetragonal and cubic polymorphs, though the details of stabilization mechanisms remain the subject of ongoing research interest [27-31]. Especially in the case of highly substituted multi-component solid solutions, a priori mapping of the stabilized phase field and prediction of lattice parameters is nigh impossible. Therefore it can be difficult to distinguish cubic from tetragonal hafnia solid solutions for compositions displaying low tetragonality, or tetragonal hafnia solid solutions from related orthorhombic compounds (e.g., HfTiO4) for phases with similar degrees of lattice distortion. In the following descriptions, phased identified as cubic stabilized defect fluorite-type hafnia solid solutions or orthorhombic HfTiO4 can thus be taken to be synonymous with hypothetically indistinguishable tetragonal polymorphs. Moving from the front (laser exposed) surface inward, the four observed scale layers displayed maximum thicknesses of 600 lm, 220 lm, 450 lm, and 90 lm, respectively, for a maximum total  \\x0c', 'D.W. Lipke et al. / Corrosion Science 80 (2014) 402-407  405  Fig. 3. (a) Cross-sectional backscattered electron micrograph of specimen B2 post-oxidation (2210 °C for 1500 s under 90 m/s airflow) exposing four distinct product layers. Arrows indicate the endpoints of thin transverse cracks thought to have formed upon cooldown and arresting within layer (1); (b) EBSD orientational maps for cubic, monoclinic, and orthorhombic indexed phases, respectively; (c) XRD patterns taken from destructively sampled layers with major peaks belonging to the {1 1 0}, {1 1 1}, and {2 0 0} symmetry-related planes in monoclinic, tetragonal, orthorhombic, and/or cubic systems; (d) WDS elemental mapping in the coarse-grained multi-phase region near the transition from layer (2)ext. to layer (3); and (e) WDS elemental line profiles for Cr and Ti across the same region mapped in (d).  Fig. 4. (a) EBSD orientational maps for cubic, monoclinic, and orthorhombic indexed phases, respectively, in the viscinity of the oxide-carbide interlayer, along with a representative backscattered electron micrograph from a nearby region for reference only; (b) WDS elemental line profiles for C and O across the oxide-carbide interlayer; and (c) Raman spectrum taken from within the oxide-carbide interlayer displaying active modes associated with nanocrystalline graphite.  two active modes. Peak characteristics, such as Raman shift, full width half maxima (FWHM), and relative intensity ratio of these modes are consistent with literature values for the occurrence of disordered nanocrystalline graphite [32]; however, no other Raman active modes could be detected. The precise character of the oxide-carbide interlayer remains ambiguous. The presence of nanocrystalline graphite accounts for at least a portion of the carbon content, although the form and amount of carbon incorporated into the oxide-based matrix is unknown. Further, the oxide-based matrix may be amorphous or  possibly nanocrystalline with crystallite sizes below the detection threshold of EBSD instrumentation. Additional characterization studies using transmission electron microscopy and/or diffraction techniques with high spatial resolution are suggested to resolve this uncertainty. Having described the complex structure of the oxidation scale, we may now remark on some of the consequences of the compositional design of the specimen. Perhaps most evident is the loss of oxidation product refractoriness of the batch composition: while pure hafnia melts at 2820 °C, the onset of melting of the tested  \\x0c', '406  D.W. Lipke et al. / Corrosion Science 80 (2014) 402-407  specimen seems to have occurred between 2210 °C and 2350 °C. Unfortunately, this might be an unavoidable consequence of alloying hafnia with substantial amounts of secondary phases, with perhaps the sole exception being rare-earth elements [1,11]. Second, the intrinsically dynamic environment experienced at ultra-high temperatures allows for thermally-activated solid state diffusion processes to proceed without kinetic inhibition, as especially evidenced by the non-uniform distributions of chromium and titanium across the relatively dense portion of the inner scale. Indeed, in the absence of differential cationic diffusion rates HfTiO4 would not have been expected to form based upon the elemental composition of the starting carbide composite. This mass redistribution effect illustrates a potentially important design concept for materials scientists working on compositional development of ultra-high temperature oxidation-resistant materials: relatively minor solid solution (or, perhaps, second phase) additions may promote the formation of pore-filling liquids locally enriched in outward-diffusing species that might serve as secondary oxygen barriers and/or in situ sintering agents. The low level of apparent porosity beneath the highly porous outer region would be consistent with the presence of a pore-filling liquid phase. Examination of the HfO2-TiO2 phase diagram (see Fig. 5) indicates that the observed multi-phase mixture consisting of HfTiO4 and hafnia would have been expected to be at least partially liquid at the average test temperature of 2210 °C [33]. However, in the absence of higher order phase diagrams including yttria, tantala, and chromia constituents, no definitive conclusion can be drawn as to whether this region developed a ‘porous framework/liquid infill’ morphology during testing similar to that responsible for self-healing behavior in HfB2/SiC composites. The oxidative performance of the tested composition bears superficial resemblance to previously reported recession rates of pure hafnium carbide and hafnium carbide-tantalum carbide composites [15,16,21]. However, we believe the testing regime employed in this study is more representative of material behavior in practice than commonly reported techniques for three primary reasons. First, the high shear velocity imposed by wind tunnel airflow can rapidly remove unsupported liquids and volatile species from the surface. Stagnant testing by comparison would be  expected to underestimate ablation rates. Second, laser illumination allows for better control over the local reactive atmosphere than does the popular oxyacetylene torch test method. Third, extended duration testing allowed for long-term oxidation behavior to be evaluated. Given the important role of solid state diffusion in establishing several of the observed phases in this study, this behavior might not have been reproduced in shorter duration tests. Further compositional development of hafnium carbide-based ceramic composites for ultra-high temperature oxidation protection is recommended. Strategies for identifying suitable substitutional solid solution or second phase additions that limit the overall loss of refractoriness of oxidation products while promoting in situ formation of a secondary oxidation barrier should be explored. Although this study did not conclusively demonstrate that the oxide-carbide interlayer was the primary impediment to oxygen ingress, this interlayer nonetheless represents one of the most refractory amorphous oxides known or, at the very least, comprises a phase capable of retaining nanocrystallinity even after extended exposure to ultra-high temperatures (i.e., it is apparently grain growth inhibited). Whichever the case may be, this interlayer merits continued study with regards to the effects of chemical composition on atomic structure, mechanism of formation, and oxygen permeability.  4. Conclusions  Chromium carbide has been demonstrated to be a useful transient liquid sintering agent for the synthesis of ultra-high temperature hafnium carbide-based solid solution ceramic composites. Consolidated samples of two batch compositions were subjected to oxidation testing at temperatures exceeding 2000 °C. The ‘A’ (Ta-rich) batch composition displayed compromised refractoriness due to excessive liquid formation upon oxidation at 2190 °C. The ‘B’ batch composition showed significant recession due to melting after experiencing a peak surface temperature of 2350 °C, yet exhibited a complex, adherent multilayer oxide scale after 25 min exposure at 2210 °C under continuous airflow of 90 m/s. Due to difficulties in phase discrimination stemming from uncertainty of the degree of tetragonality in highly substituted hafnia solid solutions, distinguishing cubic stabilized defect fluorite-type hafnia or orthorhombic HfTiO4 from tetragonal stabilized hafnia solid solutions with similar lattice distortions is nearly impossible. Though the presence of the latter hypothetical phase cannot be ruled out, the scale apparently consisted of four distinct layers: (1) 90 lm thick oxygenand carbon-containing interlayer featuring nanocrystalline graphite within an amorphous or nanocrystalline oxide matrix, (2) 450 lm thick multi-phase layer with low apparent porosity comprising orthorhombic HfTiO4, cubic stabilized defect fluorite-type hafnia, and nanocrystalline monoclinic hafnia (3) 220 lm with graded grain size distributions and proportions, thick highly porous monoclinic hafnia, and (4) 600 lm thick outermost highly porous cubic stabilized defect fluorite-type hafnia. The HfTiO4-hafnia phase mixture observed in layer (2) would have been expected to be at least partially liquid at test temperature, presumably filling the prior pore volume and inhibiting rapid oxygen ingress. If so, this ‘porous framework/liquid infill’ morphology shows promise as a design strategy for a secondary oxygen barrier, creating the low oxygen partial pressure conditions necessary for the formation of the oxide-carbide interlayer generally regarded as the foremost oxidation-resistant composition.  Acknowledgements  Fig. 5. HfO2-TiO2 phase diagram. Reprinted with permission from [33].  This research was performed while the author (D.W.L.) held a National Research Council Associateship Award at the Air Force Re \\x0c', 'D.W. Lipke et al. / Corrosion Science 80 (2014) 402-407  407  search Laboratory, Edwards AFB, CA 93524 USA. We gratefully acknowledge the assistance of Ron Witt (EBSD Analytical Inc.), Marietta Fernandez (AFRL/RQRC), Nick Botto (UC Davis), Alan Hicklin (UC Davis), Dan Siebert (LHMEL), and John Bagford (LHMEL) for their expertise and assistance with material testing and characterization efforts. Work at UC Davis was funded by the Office of Naval Research under award N00014-12-1-0196.  Appendix A. Supplementary material  Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.corsci.2013 .11.049.  References  [3]  graphite, (October P a r t IV  [1] High-Temperature Oxidation-Resistant Coatings: Coatings for protection from oxidation of superalloys, refractory metals, and graphite. National Academy of Sciences/National Academy of Engineering 1970. ISBN 0-309-01769-6. [2] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000C+ hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887-5904. J.M. Criscione, et al., High temperature protective coatings for ML-TDR-64-173 , Part I (December 1964) AD604463 , Part II 1 9 6 4 ) AD 6 0 8 0 9 2 , P a r t I I I (D e c em b e r 1 9 6 5 ) AD 4 7 9 1 3 1 , (November 1966) AD805438. [4] A.H. Heuer, V.L.K. 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Ceram. 35 (7-8) (1996) 484-496. translated from Porosh. Metall. 7-8 (1996), 182-196. [27] S. Raghavan, H. Wang, R.B. Dinwiddie, W.D. Porter, R. Vaben, D. Stover, M.J. Mayo, Ta2O5/Nb2O5 and Y2O3 Co-doped zirconias for thermal barrier coatings, J. Am. Ceram. Soc. 87 (3) (2004) 431-437. [28] P. Li, I.W. Chen, J.E. Penner-Hahn, Effect of dopants on zirconia stabilization - an X-ray absorption study: I, trivalent dopants, J. Am. Ceram. Soc. 77 (1) (1994) 118-128. [29] P. Li, I.W. Chen, J.E. Penner-Hahn, Effect of dopants on zirconia stabilization - an X-ray absorption study: II, tetravalent dopants, J. Am. Ceram. Soc. 77 (5) (1994) 1281-1288. [30] P. Li, I.W. Chen, J.E. Penner-Hahn, Effect of dopants on zirconia stabilization - an X-ray absorption study: III, charge-compensating dopants, J. Am. Ceram. Soc. 77 (5) (1994) 1289-1295. [31] A. Navrotsky, Thermodynamics of solid electrolytes and related oxide ceramics based on the fluorite structure, J. Mater. Chem. 20 (2010) 10577-10587. [32] A.C. Ferrari, J. Robertson, Interpretation of Raman spectra of disordered and amorphous carbon, Phys. Rev. B 61 (20) (2000) 14095-14107. J.P. Coutures, J. Coutures, The system HfO2-TiO2, J. Am. Ceram. Soc. 70 (6) (1987) 383-387.  [33]  \\x0c']"
},{
  "_id": 277,
  "PDF": "Ultra-High-Temperature Ceramic Coatings for Oxidation Protection of Carbon–Carbon Composites.pdf",
  "Text": "['Ultra-High-Temperature Ceramic Coatings for Oxidation Protection of  Carbon-Carbon Composites  Erica L. Corral*,w  and Ronald E. Loehman**  Sandia National Laboratories, Ceramic Processing and Inorganic Materials Department, Albuquerque, New Mexico  87106  Carbon-carbon (C-C) composites are attractive materials  for  hypersonic flight vehicles but they oxidize in air at temperatures 45001C and need thermal protection systems to survive aerothermal heating. We investigated using multilayers of high temperature ceramics  such as ZrB2 and SiC to protect C-C against oxidation. Our approach combines pretreatment and  processing  steps  to  create  continuous  and  adherent  high temperature ceramic coatings from infiltrated preceramic poly mers. We tested our protective coatings at temperatures above 26001C at controlled cold-wall heat flux profiles reaching a maximum of 680 W/cm2.  the National Solar Thermal Testing Facility using  I.  Introduction  ADVANCED thermal protection systems (TPS) are needed to mitigate effects of aerothermal heating that otherwise limit the performance of hypersonic flight vehicles. A variety of mis sion profiles are being proposed with different thermal loads and  temperatures. For example, a single use, unmanned vehicle that  has a flight profile lasting only minutes can reach a maximum temperature of B28001C.1 In order to develop these hypersonic  vehicles, new materials are needed to mitigate effects of aerothermal heating temperatures in excess of 28001C, while main taining tensile strength and oxidation resistance in these extreme  environments. C-C composites are attractive materials for hy personic flight vehicles but they oxidize in air at temperatures 45001C and need TPS to survive aerothermal heating.  (1)  Desired Atributes of a Single Use, Oxidation Protective  Ceramic Coating for C-C Composites  It is widely known that for ceramic coatings to provide oxidation protection for C-C composites they should be2,3:  (1)  adherent and continuous to provide erosion resistance  and limit the mechanisms of evaporation;  (2)  prevent oxygen permeation through the coating and re duce the diffusion of oxygen to the substrate;  (3)  mechanically compatible to resist coating spallation due  to thermal expansion mismatch during extreme thermal heating;  (4)  easily processed so that they can be made reproducibly  and reliably.  An adherent and continuous coating should provide erosion  resistance and stay intact. Thus the outer coating layer should  not be a material that oxidizes readily. If such a material must be  used because of  some other design constraints,  then an addi tional engineered mechanism should be introduced that  limits  oxidation and provides  integrity to the porous external oxide  coating. For example, a multilayer coating configuration that  has a boron-rich bottom layer and a SiC outside layer is pro tective if, on heating the boron-rich layer oxidizes  to form a  sealing glass that penetrates through the cracks in the SiC layer and limits oxidation of the coating.4  A coating that prevents oxygen permeation should be free of  defects, continuous, and have a low rate of oxygen diffusion.  This would suggest  that  there should be a low rate of oxygen  diffusion through the outer coating material or any subsequent  oxidation product.  A coating that is mechanically compatible with the substrate  during extreme thermal heating should be used as the inner lay er. If the coating material  is not mechanically compatible with  the C-C composites then coating spallation and stress-induced  cracking will result in coating defects. Methods to promote ad hesion of dense coatings and prevent delamination should con sider materials  that have  favorable wetting properties on the  C-C composite. Materials that are candidates for inner layers  are converted carbides and borate glasses. The latter combine  thermal stability with appropriate viscosity and wetting to pro vide protection over a wide range of temperatures, both as coatings and as constituents of the carbon body.3 Special atten tion should be given to the thicknesses of  the coating layers in  order to limit cracking and the quantity of oxidation products.  Lastly, pretreatments  to the C-C composite are important  in  order to promote better adhesion of the coating material during  processing.  Possible processing methods for developing coatings with fine  microstructures include single or multiple methods such a sputter deposition,4 electroplating,4 electron beam irradiation,5 liquid precursor methods,6,7 slurry coating,8 pack cementation,9 and chemical vapor deposition (CVD).10,11 Each method has  specific advantages for control of coating thickness and material  composition that should lend to a wide range of coating micro structure and configurations.  (2)  UHTC Coating Processing Approach  Ultra-high-temperature ceramics (UHTCs) have been identified  as a class of materials with the potential to withstand extreme aerothermal heating environments.12,13 For example, diboride  and carbide-based ceramics possess melting temperatures in excess of 30001C. Table I lists a range of candidate high-temper ature ceramics that have decomposition or melting temperatures above 16001C.  It  should be noted that  the table lists melting  temperatures  for materials  that do not all  exhibit  congruent  melting, the temperature at which a solid substance changes to a  liquid of  identical  chemical  composition at a given pressure.  Based on our research goal of developing an oxidation protec tive coating for C-C composites that survive for 10 min above 32001C, we have chosen to work with SiC and ZrB2 to create adherent, continuous coatings for enhanced oxidation protection of these C-C composites. Prior work by others and us12-14  Y. Blum—contributing editor  This work is presented at the Air Force Office of Scientific Research Sponsored Work shop on Ultra High Temperature Ceramic Materials, Menlo Park, CA, July 23-25, 2007.  This work is  supported by the Laboratory Directed Research Development Program at  Sandia National Laboratories. Sandia is a multiprogram laboratory operated by Sandia  Corporation, a Lockheed Martin Company,  for  the U.S. Department of Energy under  Contract DE-AC04-94AL85000.  *Member, The American Ceramic Society.  **Fellow, The American Ceramic Society.  w  Author to whom correspondence should be addressed. e-mail: elcorra@sandia.gov  Manuscript No. 23768. Received September 19, 2007; approved January 18, 2008.  Journal  J. Am. Ceram. Soc., 91 [5] 1495 - 1502 (2008)  DOI: 10.1111/j.1551-2916.2008.02331.x  r 2008 The American Ceramic Society  1495  \\x0c', 'strongly suggested that no single  coating material would be  sufficiently protective. The selection of ZrB2 was based on its high melting point. We chose SiC because it could provide low er-temperature oxidation protection and the combination of SiC  and ZrB2 would synergistically provide oxidation protection over a wide range of temperatures. Thus, our approach has been  to develop multilayer  ceramic coatings engineered to provide  necessary oxidation resistance for C-C composites.  We have developed multilayer coatings of ZrB2 and SiC on C-C composites. We used a preceramic polymer/ceramic slurry  infiltration to coat and infiltrate the C-C composites. In order to  make reliable UHTC coatings and to optimize coating infiltra tion and adhesion on C-C composites we had to thoroughly  understand the physical and chemical properties of the material  to be coated, which is why composite processing variations and  pretreatments were performed for C-C composites.  (3)  C-C Composite Pretreatment and Motivation  Because C-C composites are the structural component of  the  proposed TPS, we need to understand their  initial processing  and how it affects pore structure, surface chemistry, and bulk  crystallinity in order to apply oxidation protective ceramic coat ing materials reproducibly. The characterization of C-C com posites shows that  the processing routes and the nature of  the  fiber performs strongly influence the microstructure of the composites.15-18 Variation in the type of fibrous perform, the liquid  consolidation, and the final heat  treatment allows one to en hance drastically the graphitization of the carbon matrix and the fibers, and the size of graphite crystallites.18 The pretreatment  enhances the absorbance of ceramic coatings by controlling the  physical and chemical properties that affect coating infiltration and adhesion.19  II.  Experimental Procedure  (1)  C-C Composite Materials and Processing  The C-C composite supplier (Hitco Carbon Composites, Gar denia, CA) performed multiple variations of C-C processing  and postprocessing heat treatments at our request. A list of the  processing conditions and properties measured at Hitco for the  samples they supplied are shown in Table II. The C-C compos ite discs were a 2D carbon-carbon composite comprising T300satin weave, densified to B1.5-1.6 3K fibers in an 8-harness g/cm3 by chemical vapor infiltration (CVI). Samples were cut at  an angle to the plies to simulate a shingle lay-up. The processing  variables and purpose for pretreatment are as follows: treatment at 16001C for 4 h under  (1)  postprocessing heat  vacuum to modify adhesion between the  composite and the  coating. The  heat  treatment  partially  removes  the  surface  chemical groups and cleans  the  carbon fiber  surfaces by de functionalizing them and enhances crystallinity and removes any  amorphous  carbon from the  surface. This  creates a uniform  carbon surface composition and with a higher degree of graph itization, which may allow better wetting and infiltration of the  liquid precursor  coating. The heat  treatment also affects  the  pore structure between the fibers and provides larger channels  for UHTC precursor infiltration into the substrate. This is due  to the  increase  in graphitization of  the fibers  from the heat  treatment, resulting in more ordered and tightly packed fibers  that create larger void spaces between them;  (2)  CVI times of 150 and 50 h. The CVI process is used to  modify final densities of  the C-C composites through internal  deposition of graphitic carbon. The standard high-density CVI  process is 150 h, whereas the low-density composites are infil trated for 50 h. The low-density composites have surfaces with a  uniform porous  structure  that  allow for maximum coating  infiltration.  (2)  UHTC Coating Materials and Processing  The coating materials evaluated in this  investigation were in tended to serve as  the outer  layers of  the TPS. A number of  Table II.  C-C composite Physical Properties before Ceramic  Coating Infiltration  C1109  C1110  C1113  C1114  Filler type  C-black  C-black  C-black  C-black  CVD time (h)  50 16001C  150 16001C  50  150  Heat treatment Final density (g/cm3)  None  None  1.5  1.6  1.5  1.6  Total voids  21.5%  16.7%  21.7%  15.9%  C-C, carbon-carbon; CVD, chemical vapor deposition.  Table I.  UHTCs for use as Potential Thermal Protective  Materials  Ceramic  material  Decomposition & melting  temperature (Td and Tm)  Density  (g/cm3)  Oxidation products  (in air)  SiC22  Td 5 16001C Tm 5 27301C Tm 5 32001C Tm 5 32501C Tm 5 38901C Tm 5 35321C  3.22  SiO21CO2  ZrB2 HfB2 HfC28 ZrC28  21  6.08  ZrO21B2O3 HfO1B2O3 HfO1CO2 ZrO21CO2  27  10.50  12.22  6.73  UHTC, Ultra-high-temperature ceramics.  Fig. 1. Mercury intrusion porosimetry pore size distribution for carbon-carbon (C-C) composites (a) C1109 and C1113, and (b) C1110 and C1114 with  and without heat treatment, respectively.  1496  Journal of the American Ceramic Society—Corral and Loehman  Vol. 91, No. 5  \\x0c', 'May 2008  UHTC Coatings for Oxidation Protection of C-C Composites  1497  Fig. 2.  (a) A scanning electron micrograph (SEM) of the surface of a carbon-carbon (C-C) composite coated several times using only the precursor to  SiC. The coating is discontinuous and does not provide for oxidation protection to the exposed areas of C-C composite. SEM of C-C composite top  surfaces after coating with the SiC slurry/precursor mixture (b) two times (c) covering 89% of the composite area and (d) three times (e) covering 97% of  the composite area. The red color indicates the coating area used to calculate the average area covered. SEM cross-sections for coated C-C composites  after (f) one infiltration cycle and (g) after three infiltration cycles.  precursors  to SiC were  evaluated  as  infiltrated coatings on  C-C composites before selecting a commercially available pre also synthesized using published methods6 for use as ZrB2 coatings. Ceramic slurries of SiC particles suspended in the liquid  cursor (SiC SMP-10, StarFire Systems, Malta, NY). This ally precursor to SiC were also used to build thicker and more con lhydridopolycarbosilane precursor to SiC was available in large  tinuous coatings. The  commercially available SiC slurry (SiC  quantities and could be handled in air. Precursors to ZrB2 were  Matrix Slurry MS-10, StarFire Systems, Malta, NY) was mixed  \\x0c', 'with the SiC SMP-10 precursor to optimize the slurry/precursor  viscosity for maximum infiltration into the C-C composites. We  used an 80 wt% slurry-20 wt% precursor mixture. The dip coat  and pyrolysis cycle was  repeated up to three times to make a  continuous coating. SiC and ZrB2 coatings were applied as single layers and processed using vacuum pressure-assisted dip coating and infiltration, followed by a pyrolysis heat treatment under flowing Ar (21C/min to 4501C, hold for 4 h; 41C/min to 11001C, hold for 1 h; 101C/min to 1001C; furnace cool).  (3)  Solar Furnace Testing: Thermal Shock at High Temperature and Severe Heat Flux Conditions  Low cost,  reliable  testing methods  for  evaluating oxidation  protection materials are needed to advance TPS materials de velopment and gain an understanding of coating materials per formance at temperature. Our multilayer coatings were tested at (1700-26001C) using multiple facilities  high temperatures  that  allow for quick testing and processing iterations during coating  material development. The Solar Furnace Facility was used to  test coatings for thermal shock at high temperature and severe  heat-flux conditions. It has the ability to accurately control a variable heat-flux up to 800 W/cm2 while continuously monitoring the specimen surface temperature up to 26001C. This test  is the most relevant for heat-flux, high temperature, and expo sure time to a related hypersonic material application. Because  tests were conducted in subsonic air-flow, they do not exhibit all  aspects of the application.  The Solar Furnace Facility uses a heliostat that tracks the sun  to direct  sunlight onto a mirrored parabolic dish. The  focal  point of the dish does not move, which simplifies sample place ment. The power level of the furnace is adjusted using an atten uator  that works  like a Venetian blind located between the  heliostat and the dish. The furnace provides 16 kW total thermal  power with time-dependent control of the thermal flux to a maximum of 800 W/cm2. The heat-flux is adjusted by a com puter-controlled  attenuator  that permits  a  time-varying flux  profile. Samples measuring 1.59 cm in diameter and 0.64 cm  thick were heated under a time varying profile to study the per formance of the coating treatment. Specimen surface temperatures were monitored with optical pyrometers at 5.2 and 1.7 mm  wavelengths and the test specimen surface conditions were ob served with video and infrared cameras. The surface temperabe o1001C. tested separately and the duration of each test  ture  uncertainty was measured  to  Each  test  specimen was  was dictated by the heat-flux profile. The tests were open to air  and the specimen was exposed to shear airflow, approximately  25 m/s. The maximum heat flux reached at testing is dependent  upon cloud cover, which restricted testing even when the sky  was only partially cloudy. For optimum weather test conditions, is B25% but, of  the  relative  humidity  course,  the  relative  humidity at the testing temperature is much lower.  (4) Materials Characterization  Mercury intrusion porosimetry (MIP) was used to characterize  the porosity of  the different C-C composites before  ceramic  coating infiltration on C-C composite discs (Autopore IV 9500,  Micromeretics  Instrument Corp., Norcross, GA). The phase  compositions of the ceramic coatings and uncoated C-C composites were analyzed with an X-ray diffractometer (CuKa ra diation, X’Pert PRO, PANalytical, Almelo,  the Netherlands)  and scanning electron microscope (SEM, Zeiss SUPRA 55VP  FE-SE, Carl Zeiss SMT Inc., Peabody, MA) equipped with an  energy-dispersive X-ray microanalyzer  (EDX, Rontec  SDD  quad detector, XFLASH Quad4040, Bruker AXS Microanaly sis, Berlin, Germany). X-ray photoelectron spectroscopy (XPS,  AXIS Ultra DLD, Kratos Analytical, Columbia, MD) was used  to characterize the surface chemistry of C-C composites before  coating infiltration.  Polished sections of the uncoated and coated C-C composites  were ground flat using 600 grit SiC papers then prepared with to a 1.0 mm finish. Final  successively finer diamond abrasives  polishing was achieved using a two-step process, starting with 0.3 mm alumina, then 0.04 mm silicon dioxide using vibratory  polishers.  III.  Results  (1)  C-C Physical Properties and Surface Chemistry  Characterization of  the C-C composite pore structure, surface  chemistry, and crystallinity are necessary in order to apply coat ing materials  reproducibly. The porous  structure of  the C-C  composite determines the adsorption capacity of coating mate rial. Figs. 1(a) and (b) show the MIP pore size distribution for  C-C composites C1109, C1113, C1114, and C1110. Heat treat ment does not change the pore size distribution for C-C com posites  that  have  been  densified  for  only  50  h,  as  seen  in  Fig. 1(a). However, heat treatment does cause the pore size dis tribution of the 150 h densified composites to shift to lower pore  sizes, as seen in Fig. 1(b). As expected, materials densified using  the standard CVD cycle time (150 h) have more deposited amor phous carbon and thus experience more graphitization during  the postprocessing heat  treatment. The heat  treatment graph itizes  (crystallizes)  the amorphous carbon, resulting in a more  ordered crystallite  structure  and alters  the porous  structure.  Fig. 3.  (a) A scanning electron micrograph micrograph of the carbon-  carbon (C-C) composite surface coated three times using the precursor  to ZrB2 and SiC mixture. (b) The SiC/ZrB2 multilayer coating covered 99% of C-C composite C1111. (c) The energy-dispersive X-ray spect rometry elemental map shows  the areas  covered with Zr  (green), C  (blue), and Si (red).  1498  Journal of the American Ceramic Society—Corral and Loehman  Vol. 91, No. 5  \\x0c', 'May 2008  UHTC Coatings for Oxidation Protection of C-C Composites  1499  Fig. 4.  Scanning electron micrograph (SEM) micrographs, energy-dispersive X-ray spectrometry (EDS) elemental spot analysis and X-ray diffraction  spectra for the SiC/ZrB2 coating that was unprotective during solar testing. (a) SEM micrographs of the tested surface show oxidation of the C fibers and (b) small SiC particles between the oxidized carbon-carbon weaves. (c) The EDS elemental analysis of the particles shows they contain Si. (d) Magnified  image of SiC particles that were analyzed for Si using EDS elemental analysis.  Thus for the 150 h densified composites, heat treatment increas heat treatment. As observed by both XPS and XRD, heat treat es  the  overall  accessible  pore  volume  porosity  of  the C-C  ment  enhances graphitization of  the C-C composites, giving  composites.  larger crystallites  that are better oriented and can be used to  The effect of heat treatment on the surface chemistry of C-C  alter the pore size distribution within the C-C composite.  composites was analyzed using XPS. The C 1s peaks are nar rower after heat  treatment and have lower  full width at half  maximum (FWHM) values,  indicating a more graphitic char (2)  Processing Preceramic Polymers and Ceramic Slurry  acter after heat  treatment. The graphite crystals are better ori Coatings  ented and increase  in size with heat  treatment. The higher  FWHM value without heat treatment also suggests that the un treated carbon surface has more defects due to the presence of  functional groups on the surface than after heat treatment and is less ordered.20 The functional groups present on the surface of  the C fibers  interact with adsorbates. Thus heat  treatment re moves  surface  functional groups and improves  the  structural  organization of the C-C composite. The bulk graphitic charac ter of  the C-C composite was also characterized using X-ray  diffraction (XRD). We found that the major graphitic (002) reflection at 2y 5 241  in intensity with postprocessing  increases  Fig. 5.  The X-ray diffraction spectrum shows  low graphite intensity  reflections and small SiC peaks  for a SiC/ZrB2 provide oxidation protection to the carbon-carbon composite.  coating that did not  Using particle filled precursor mixtures allows us  to back-fill  cracks and pores that result after coating with only preceramic  polymers, as seen in Fig. 2(a). The figure shows the surface of a  C-C composite after several coatings and heat treatments using  only the precursor to SiC. As one can see, there are large gaps  between the crystallized SiC regions that leave large unprotected  areas of C-C composite. The best SiC coating composition was  found to be a mixture of 80:20 (wt%) Matrix Slurry MS-10:  SMP-10 precursor. Figure 2(b)  shows  the  surface of a C-C  composite coated two times with the SiC slurry/precursor mix ture. Figure 2(c) shows that the SiC coating covered up to 89%  of the area after only two infiltration and coating cycles. Figure  2(d) shows the surface after coating three times using the SiC  mixture. Figure 2(e) shows that the SiC coating covers up to 97%  of the area after three coating cycles. Figure 2(f) shows a polished  cross-section of a coated C-C composite after only one dip coat  cycle and Fig. 2(g) shows the cross-section of a C-C composite  coated three times. These figures show that the viscosity of our  SiC coating mixture was appropriate  to infiltrate  small pore  channels within the  composites after only one  coating cycle.  Further, use of multiple coating cycles shows we can make coat ings with the SiC mixture that are uniform, adherent, and continuous. The average coating thickness observed was 20-35 mm.  We also used a precursor to ZrB2 to infiltrate and coat C-C composites to build multilayer coatings. Figure 3 shows a SEM  micrograph of a C-C composite surface with SiC and ZrB2 precursors as coated as multilayers. The surface of the coated C-C  composite exhibits dark gray regions that are identified as SiC  with a finer microstructure that is the ZrB2 coating. Figure 3(b)  \\x0c', '1500  Journal of the American Ceramic Society—Corral and Loehman  Vol. 91, No. 5  shows that the average area fraction coated on the surface of the  agglomerated SiC particles that survived the test embedded be C-C composite was greater  than 99%. Figure 3(c)  shows  the  tween them. The EDS elemental analysis of  the particles con energy-dispersive X-ray spectrometry (EDS) elemental map for  firmed that they contain Si, as seen in Figs. 4(c) and (d). Figure 5  Si (red), Zr (green), and C (blue) from the coated area. The more  shows the XRD spectrum of the tested specimen surface, where  uniformly coated regions are from the SiC mixture, whereas the  intensities of the expected graphite peaks are low due to the ox ZrB2 regions were finer and less continuous due to the use of a nonparticle-filled liquid precursor coating. These micrographs  suggest that the precursors to ZrB2 and SiC did not uniformly adhere to the C-C composite as layers. This observation further  illustrates  the  complexity  of  reproducibly making  uniform  coatings using liquid preceramic polymers.  (3)  Analysis of Multilayer Coatings on C-C After High  Temperature Testing  The following solar  test  results  represent  the most  significant  and promising findings  for oxidation protection of C-C com posites using ceramic coatings for application in TPS. Results  after Solar Furnace  testing are most  relevant  to the  thermal  profiles of projected flight environments as they directly dupli cate  the  severe  expected heat-flux and high temperatures  for  short  periods  of  time.  SEM micrographs,  EDS  elemental  spot  analysis,  and XRD spectra were  used  to  characterize  the coatings that were not protective during solar testing (680 W/cm242600 1C). Figures 4(a) and (b) show SEM micro graphs for the C-C composite that was not protected from ox idation. There are exposed weaves of oxidized C-C fibers with  idation of the C-C. The low intensities for the SiC peaks are due  to the failure and depletion of the coating during solar testing.  Figures 6(a) and (b) show SEM micrographs of the surface of  a SiC/ZrB2 multilayer ceramic-coated C-C composite that was protective in solar testing (680 W/cm2 for 15 s). The micro graphs show a C-C composite surface that  is coated with SiC  and decorated with islands that contain boron and are presumed  to be B2O3 particles. Figure 6(c) shows spot elemental analysis of a representative test area where points one, two, and three  were analyzed. The  small particles  contain boron, as  seen in  Figs. 6(e) and (f), and the surrounding area is covered with Si, as  seen in Fig. 6(d). The XRD spectrum, shown in Fig. 7 reveals  intense graphite reflections as well as peaks for SiC, B2O3, and SiO2, indicating that the multilayer coating protected the C-C composite  from oxidation. Figure 8 presents  electron micro graphs of  cross-sections of  samples  shown in Figs. 4 and 6,  where the coatings are (a) unprotective and (b) protective. The  cross-sections provide additional detail on how the UHTC coat ings provide oxidation protection for the C-C composite. The  cross-section in Fig. 8(a) shows that  the heated surface of  the  C-C composite has partially oxidized. Therefore the coating was  compromised and also allowed partial oxidation of  the C-C  Fig. 6.  Scanning electron micrograph (SEM) micrographs and energy-dispersive X-ray spectrometry (EDS) elemental spot analysis for a SiC/ZrB2 coating that did provide oxidation protection for the carbon-carbon (C-C) composite during solar testing. SEM micrographs (a) and (b) of the tested  surface show SiC and crystallites identified as B2O3 on the surface of the C-C composite. (c) The EDS elemental analysis of the analyzed area show the points analyzed for (d) Si and (e-f) for (b).  \\x0c', 'substrate. The partially oxidized C-C composite still contains a  small amount of SiC deeper within the composite that was not  oxidized furing the test. Figure 8(b) shows that the UHTC coat ing did oxidize near  the  test  surface. Furthermore, a greater  amount of SiC coating is still contained within the composite, as  can be seen as  the light-phase material  in the micrograph. As  expected,  the  thickness of  the multilayer  coating on the  test  surface of  the C-C composite was depleted after high-temper ature solar furnace testing.  Our results clearly show that multilayer coatings with SiC and  ZrB2 provide oxidation protection for C-C composites under severe heat-flux and high-temperature conditions. As mentioned  before, numerous tests showed that single continuous coatings  of  either SiC or ZrB2 did not protect C-C composites oxidation. Thus, these experimental results show that protective  from  coatings must be continuous and contain both SiC and ZrB2.  IV.  Discussion  (1)  High-Temperature Oxidation Mechanisms of ZrB2 and SiC  During oxidation, ZrB2 forms stable ZrO2 and B2O3, which is volatile at temperatures above 9001C. As a result, the growth  of the reaction products is governed by the rates of formation of  ZrO2 and volatilization of B2O3. The atmospheric oxidation  of ZrB2 may be described by the following reaction21:  ZrB2 þ 5=2O2 , ZrO2 þ B2O3  At high temperatures, SiC exhibits two types of oxidation be havior, depending on the ambient oxygen potential. At high  oxygen pressures, passive oxidation occurs wherein a protective film of SiO2(s) is formed on the surface by the reaction22:  2SiCðsÞ þ 3O2 ðgÞ ! 2SiO2 ðsÞ þ 2COðgÞ  At  low oxygen potentials, active oxidation occurs due to the  formation of gaseous products according to reactions:  SiCðsÞ þ 2SiO2 ðsÞ ! 3SiOðgÞ þ COðgÞ SiCðsÞ þ O2 ðgÞ ! SiOðgÞ þ COðgÞ  Since the early 1970s, there have been extensive efforts to in crease  the  oxidation  resistance  of  diborides. These  studies  showed that the addition of SiC to ZrB2 greatly increases the resistance to oxidation at high temperature.14,23-25 The thermo dynamic calculations for the oxidation of ZrB2-SiC composites by Fahrenholtz26 show that the formation of a SiC-depleted  layer beneath the outer oxide layers of ZrO2 and SiO2 is thermodynamically favored and that the addition of SiC slows down  the oxidation kinetics. These results show that no single material  composition will provide oxidation protection over a wide range  of temperatures, which is consistent with our approach to syn thesizes and test adherent continuous coatings of ZrB2 and SiC to provide for single-use, high-temperature oxidation protection  of C-C composites.  Based on our observations, coatings that are not continuous  and adherent allow the C-C composite to oxidize at high tem peratures and thus it is not protected. Such discontinuous mul tilayer coatings fail because the ZrB2 oxidizes to ZrO2 and B2O3, allowing the latter to volatilize at T4900 1C, leaving behind  ZrO2. The SiC coating also begins to oxidize to SiO2 but mains mostly SiC between the layers of C-C sheets where  re conditions are  locally very reducing. However, a continuous  SiC/ZrB2 glass coating on the  coating  is protective  and seems  to form a borate  surface  that upon cool-down contains  embedded borate crystallites  in the coating. The surrounding  majority of C-C composite is still covered with SiC after testing.  Mismatch of  thermal expansion coefficients of composite and  coatings is not an issue at high temperatures because any small  cracks are filled with the viscous liquid formed from oxidation of  ZrB2 and SiC.  V.  Conclusions and Summary  We have identified multilayer coatings that protect C-C composites from oxidation at 680 W/cm2 (Tmax426001C). The C-C composites showed oxidation resistance for tests at high heat fluxes and high temperatures at the Solar Furnace Facility. We  have produced a coating system that provides thermal protec tion for C-C composites that are exposed to extreme conditions  for  short periods of  time. Coatings with both ZrB2 and infiltrated slurries of SiC powder in liquid SiC precursor were re quired for best oxidation resistance. We found it  important  to  engineer the C-C composite surface chemistry and microstruc ture in order  to form adherent  interfaces between the coating  and the composite. Our results also suggest  that multimaterial  coatings are required for  full protection and that  the coating  compositions  and microstructures need to be  carefully  engi neered for the particular use conditions.  Acknowledgments  The authors thank Marlene Chavez of Sandia National Laboratories for pro cessing the coatings and Mike Edgar and Tim Moss for solar furnace testing at the  National Solar Thermal Testing Facility. Materials characterization performed by  Denise Bencoe, Tom Hinklin, James A. Ohlhausen, William Wallace, Alice Kilgo,  Fig. 8.  Polished  cross-sections  of multilayer  ultra-high-temperature  ceramics coated carbon-carbon (C-C) composites after  solar  furnace  testing, provide additional detail on how they performed as (a) unpro tective and (b) protective coatings.  (a) The C-C composite is partially  oxidized and depleted of SiC near the surface. (b) The C-C composite is  fully protected and rich in SiC near the surface.  Fig. 7.  The X-ray diffraction spectrum shows high graphite intensity  and peaks  for SiC, B2O3, and SiO2 after solar testing for a SiC/ZrB2 coated carbon-carbon composite.  May 2008  UHTC Coatings for Oxidation Protection of C-C Composites  1501  \\x0c', '1502  Journal of the American Ceramic Society—Corral and Loehman  Vol. 91, No. 5  and Bonnie McKenzie are also appreciated. The authors are also grateful  for  the  contributions  from the TPS research team members  at Sandia National  Laboratories.  References  1S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem,  ‘‘Evaluation  of Ultra-High Temperature Ceramics  for Aeropropulsion Use,’’  J. Eur. Ceram. Soc., 22, 2757-67 (2002). 2J. E. Sheehan,  ‘‘Oxidation Protection for Carbon Fiber Composites,’’ Carbon,  27 [5] 709-15 (1989). 3J. R. Strife and J. E. Sheehan,  ‘‘Ceramic Coatings for Carbon-Carbon Com posites,’’ Am. Ceram. Soc. Bul., 67, 369-74 (1988). 4J. E. Sheehan,  ‘‘High-Temperature Coatings on Carbon Fibers and Carbon-  Carbon Composites,’’ pp. 223-66 in Carbon-Carbon Materials and Composites,  Edited by J. D. Buckley and D. D. Edie. William Andrew Inc, Norwich, 1993. 5M. D. Alvey and P. M. George,  ‘‘ZrPt3 as a High-Temperature, Reflective, Oxidation-Resistant Coating for Carbon-Carbon Composites,’’ Carbon, 29 [415]  523-30 (1991). 6J. A. Jensen, J. E. Gozum, D. M. Pollina, and G. S. Girolami,  ‘‘Titanium,  Zirconium, and Hafnium Tetrahydroborates as ‘‘Tailored’’ CVD Precursors for  Metal Diboride Thin Films,’’ J. Am. Chem. Soc., 110, 1643-4 (1988). 7Y. Yan, Z. Huang, S. Dong, and D. Jiang,  ‘‘New Route to Synthesize Ultra Fine Zirconium Diboride Powders Using Inorganic-Organic Hybrid Precursors,’’  J. Am. Ceram. Soc., 89 [11] 3585-8 (2006). 8D. W. McKee,  ‘‘Oxidation Behavior and Protection of Carbon-Carbon Com posites,’’ Carbon, 25 [4] 551-7 (1987). 9Q. G. Fu, H. J. Li, X. H. Shi, K. Z. Li, and G. D. Sun,  13M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C Hypersonic Aerosurfaces: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39, 5887-904 (2004). 14F. Monteverde and L. Scatteia,  ‘‘Resistance to Thermal Shock and to Oxida tion of Metal Diborides-SiC Ceramics for Aerospace Application,’’ J. Am. Ceram.  Soc., 90 [4] 1130-8 (2007). 15P.  Ehrburger  and  J.  Lahaye,  ‘‘Characterization  of  Carbon-Carbon  Composites-I,’’ Carbon, 19, 1-5 (1981). 16P. Ehrburger and J. Lahaye,  ‘‘Characterization of Carbon-Carbon Compos ites-II,’’ Carbon, 19, 7-10 (1981). 17M. L. Minus and S. Kumar,  ‘‘The Processing, Properties, and Structure of  Carbon Fibers,’’ J. Metals, 57 [2] 52-7 (2004). 18T. Paulmier, M. Balat-Pichelin, and D. L. Queau,  ‘‘Structural Modifications  of Carbon-Carbon Composites Under High Temperature and Ion Irradiation,’’  Appl. Surf. Sci., 243, 376-93 (2005). 19Y. Nakayama, F. Soeda, and A.  Ishitani,  ‘‘Study of Carbon Fiber Matrix  Interface,’’ Carbon, 28 [1] 21-6 (1990). 20T. Takahagi and A. Ishitani, ‘‘XPS Study on the Surface Structure of Carbon  Fibers using Chemical Modification and C1s Line Shape Analysis,’’ Carbon, 26 [3]  389-96 (1988). 21R. F. Voitovich, E. A. Pagach, and L. A. Men’shikova,  ‘‘High-Temperature  Oxidation of Zirconium Diboride,’’ Metallurgiya, 54 [6] 44-8 (1967). 22S. C. Singhal, ‘‘Oxidation Kinetics of Hot-Pressed Silicon Carbide,’’ J. Mater.  Sci., 11, 1246-53 (1976). 23A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby,  ‘‘Characterization  of Zirconium Diboride Ceramics  for Thermal  Protection  Systems,’’ Key Eng. Mater., [264-268] 493-96 (2004). 24F. Monteverde, A. Bellosi,  and S. Guiciardi,  ‘‘Processing  and Properties  ‘‘Silicon Carbide  of Zirconium Diboride-Based Composites,’’  J. Eur. Ceram. Soc.,  22,  279-88  Coatings  to Protect Carbon/Carbon Composites Against Oxidation,’’ Scripta  Mater., 52, 923-7 (2005). 10W. Cermignani, T. Paulson, and C. Onneby, ‘‘Synthesis and Characterization  of Boron-Doped Carbons,’’ Carbon, 33 [4] 367-74 (1995). 11J. Chown, R. F. Deacon, N. Singer, and A. E. S. White, ‘‘Refractory Coatings on  Graphite,’’ pp. 1-81 in Special Ceramics, Edited by P. Popper. Academic Press, 1963. 12R. Loehman, E. Corral, H. P. Dumm, and P. Kotula.  ‘‘Ultra High Temper ature Ceramics for Hypersonic Vehicle Applications’’; Technical Report No. SAND  2006-2925 (Sandia National Laboratories, 2006).  (2002). 25W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of an SiC Addition on  the Oxidation of ZrB2,’’ Ceram. Bul., 52 [8] 612-6 (1973). 26W. G. Fahrenholtz, ‘‘Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 27E.  ‘‘Oxidation of ZrB2 and HfB2 Oxygen Atmosphere,’’ Poroshk. Metall., 190 [11] 77-80 (1978). 28R. F. Voitovich and E. A. Pugach, ‘‘High-Temperature Oxidation of ZrC and  I. Golovko and R. F. Voitovich,  in an  HfC,’’ Poroshk. 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},{
  "_id": 278,
  "PDF": "Ultra-High-Temperature Ceramics Based on HfB2 – 30_ SiC Production and Basic Properties.pdf",
  "Text": "['ULTRA-HIGH-TEMPERATURE CERAMICS BASED ON HfB2 - 30% SiC:  PRODUCTION AND BASIC PROPERTIES  P. S. Sokolov,1,2 A. V. Arakcheev,1 I. L. Mikhal’chik,1 L. A. Plyasunkova,1  I. F. Georgiu,1 T. S. Frolova,1 R. A. Mironov,1 A. V. Lanin,1 A. O. Zabezhailov,1  I. Yu. Kelina,1,3 and M. Yu. Rusin1  Translated from Novye Ogneupory, No. 5, pp. 48 - 55, May 2017.  Original article submitted January 12, 2017.  Dense HfB2 - 30 vol.% SiC ceramics were obtained from commercially available powders by the hot-press ing method. Their basic physicomechanical properties were measured: the flexural strength at room tempera ture was 300 - 700 MPa, the Vickers microhardness reached 20 GPa, the critical stress intensity factor was up  to 5.9 MPa·m1/2. Thermal expansion and resistance to oxidation were measured in a wide temperature range.  Keywords: ultra-high-temperature ceramics (UHTCs), ceramics based on HfB2-SiC, hot pressing (HP), haf nium diboride, silicon carbide.  INTRODUCTION  The development of new refractory and heat-resistant ce ramic composite materials  for  the aerospace industry is an  important and urgent  task. An important  issue is the protec tion of the nose and sharp edges of the wings of hypersonic  aircraft  (HSA).  Traditional  aviation materials,  such  as  heat-resistant  alloys  based  on  titanium and  nickel,  car bon-carbon composites and ceramics based on silicon car bide operate at only up to 1650°C and are no longer able to  meet the increasing requirements. Ultra-high-temperature ce ramics (UHTC) based on zirconium and hafnium diborides is  a promising class of structural materials [1]. The high melt ing point of ZrB2 and HfB2 (>3000°C) in combination with  their phase stability, high hardness, ability to form barrier ox ide coatings at extreme temperatures make them potential  candidates  for  long-term operation  at  temperatures  above  1800°C.  The key properties  (Young’s modulus, microhardness,  thermal coefficient of linear expansion (TCLE),  thermal and  electrical conductivity) of ceramics based on ZrB2 and HfB2  are  generally  similar  [1 - 3]. At  the  same  time,  hafnium  diboride has several  important advantages. First,  the kinetic  resistance to oxidation in HfB2 is approximately 2 - 3 times  higher than that of ZrB2. Second, the melting points of HfB2  and HfO2 are approximately 130°C higher than of ZrB2 and  ZrO2, respectively.4 Third, the ionization energy and stability  in a vacuum in HfB2 is also higher than in ZrB2. And, finally,  the fundamental difference in nuclear properties — the ther mal neutron capture cross section in hafnium compounds is 3  orders  of magnitude  higher  than  that  of  zirconium com pounds.  The only obvious  shortcomings of materials based on  hafnium compounds are its relatively small prevalence and a  high dispersion in the earth’s crust. Hafnium does not  form  independent minerals and usually accompanies zirconium. In  zirconium ores,  its content  is about 2 - 4% of the amount of  zirconium. Commercially available HfB2 powders are very  expensive, and in relation to this,  there are relatively few  studies of the properties of ceramics based on it  in compari son with zirconium diboride.5 The relatively high density of  HfB2 (11.20 versus 6.12 g/cm3 in ZrB2), in the opinion of the  Refractories and Industrial Ceramics  Vol. 58, No. 3,  September, 2017  304  1083-4877/17/05803-0304 © 2017 Springer Science+Business Media New York  1  ORPE “Technologiya” named after A. G. Romashin, a State Re search Center of  the Russian Federation, Obninsk, Kaluga Re gion, Russia.  2  sokolov-petr@yandex.ru  3  keline@technologiya.ru  4  This also leads to the fact  that  the temperature of  the eutectics  HfB2-MoSi2, HfB2-SiC, HfB2-B4C, HfB2-C, HfB2-HfC is  strictly higher by 100 - 200°C than in the corresponding pairs  with zirconium diboride.  5  Over the past five years, no more than 200 research papers have  been published on the synthesis of initial powders, sintering, and  the study of the properties of HfB2-based ceramics.  DOI 10.1007/s11148-017-0101-4  \\x0c', 'authors of [2], is not a disadvantage and, on the contrary, can  be useful  for moving the center of gravity of  the HSA for ward in some aerodynamic schemes.6 The object of the pres ent work is  ceramics of  the  composition HfB2 - 30 vol.%  SiC obtained by hot pressing. Experimental data on the main  properties over a wide range of  temperatures are presented  below.  EXPERIMENTAL SECTION  Source materials  Fine dispersed hafnium boride powders (purified) of do mestic production were used. As a second component of the  ceramics, abrasive silicon carbide powders of  technical pu rity were used. The phase composition and microstructure of  the initial HfB2 powders were studied by x-ray phase analy sis (XRD), optical and electron microscopy. Express analysis  of the chemical composition of HfB2 powders was performed  by x-ray spectroscopy (semi-quantitative EDXMA) using the  Quantax Bruker energy dispersive attachment. The content  of  impurities in silicon carbide powders was determined by  quantitative chemical analysis methods.  The specific surface area of  the powders was measured  by the gas filtration method on the PSKh-9 instrument,  the  pycnometric density r p was evaluated on a helium pycno meter “AccuPyk II 1340,” the particle size distribution was  determined by the laser diffraction method on the “Fritsch  Analysette 22 Microtec plus” instrument. The measurement  was  carried  out  in  distilled water with  the  addition  of  a  surfactant  and  ultrasound  treatment. The  distribution was  calculated according to the Fraunhofer theory.  Synthesis  HfB2 powders were  screened  through  a  brass  sieve,  working fractions (<40 mm) were selected. Preliminary siev ing was not used for SiC. The powders were then mixed in a  wet  process  in  a  laboratory  planetary  ball mill  and with  grinding jars made of  silicon nitride. The  stirring time  in  isopropyl  alcohol  at 150 rpm was 24 hrs. The  charge was  then  removed  from the  grinding  jars,  dried  and  rubbed  through a brass sieve with 100 micron cells. All operations  were performed in air. Samples Nos. 1 and 2 were obtained  from HfB2 powder from KTM Ltd., Nos. 3 and 4 from HfB2  from “Ekos-Ural.” SiC from the Zaporozhye Abrasive Plant  (ZAP) was used in samples Nos. 1, 3, and from the Volzhsky  Abrasive Plant (VAP) - in samples Nos. 2, 4.  The  resulting mixtures were  placed  in  a  hot-pressing  (HP) unit of an original construction. The compacting and si multaneous  sintering were  carried out  in a graphite mold  with an internal  cross  section of 70 ´ 32 mm2,  lubricated  with  boron  nitride. The maximum pressing  pressure was  about 20 MPa,  the  temperature was 1900 - 2000°C. After  HP, mechanical processing of  the ceramic was carried out  with a diamond tool (samples Nos. 1 - 3) or a spark erosion  wire (No. 4).  Analysis and preparation of ceramic samples  The actual density r a, the open porosity P, and water ab sorption W were measured by hydrostatic weighing using  distilled water at 20°C;  the geometric method was also used  to  confirm the  results.  The  theoretical  density  (for  the  HfB2 - 30% SiC composition, r  t = 8.80 g/cm3) was  calcu lated by the rule of mixtures,  taking the density of HfB2 and  SiC to be 11.20 and 3.21 g/cm3, respectively. The content of  HfO2 and B2O3 impurities was not  taken into account  in the  calculation of r t. The relative density was taken from the ra tio r  a/r  t.  To measure the flexural  strength for a static four-point  bending s fl at  room temperature,  the ceramic was  first cut  into 3.5 ´ 4.5 ´ 47.0 mm blanks. The blanks were then pol ished along the long side to a roughness of Ra ~ 0.3 mm7 with  the  diamond  tool  (100/80 mm)  and  a  finite  beam size  of  3.0 ´ 4.0 ´ 45.0 mm.  Chamfers were  formed  on  all  the  beams to minimize the effect of stress concentration. Flex ural strength s  fl  tests were carried out according to ASTM  C1161-13  using  the  Walter + Bai  AG  LFM-50  and  LFM-100 machines. The loading speed was 1.5 mm/min, the  distance between the support elements was 40 mm, and be tween the upper punches — 20 mm. Some  samples were  pre-oxidized at 1400°C for 1 h and then tested at room tem perature. A total of 40 samples were tested. The microhard ness of HV was measured at 20°C by the Vickers method  (load from 0.05 to 3 kg, indentation time 10 sec) on polished  surfaces  (Ra ~ 0.03 mm) using an EmcoTest DuraScan 50  hardness tester. The measurements were carried out at  three  points at each load. The critical stress intensity factor KIc was  estimated by the method of  indentation along the length of  radial cracks (Palmkvist cracks) that form near the angles of  a diamond pyramid imprint at a load of 3 kg in 16 points.  The phase composition of  the ceramics was determined  on the polished surface by means of an x-ray diffractometer  DRON-6 (Cu Ka-radiation). The microstructure was studied  by optical and electron microscopy with an EDXMA attach ment on an AXIO Observer (Zeiss) optical microscope and a  EVO 40 XVP (Zeiss) scanning microscope. To characterize  the samples,  the grain size distribution was studied using the  ImageJ software package. The temperature coefficient of linear  expansion TCLE (a) was measured in accordance with GOST  10978-2014 on samples measuring 4.0 ´ 4.0 ´ 50.0 mm in  Ultra-High-Temperature Ceramics Based on HfB2 - 30% SiC  305  6  The  density  of HfB2  is  close  to  the  density  of molybdenum  (10.2 g/cm  3  ) and high-temperature alloys based on it.  7  The roughness of the treated surface of the ceramic was measured  on a profilometer TR-200 in accordance with GOST 2789.  \\x0c', 'air in the range from 20 to 1100°C on a dilatometer “Netzsch  DIL 402C,” previously calibrated using an alumina ceramic  standard of similar dimensions. The data were represented as  a mean TCLE in the range of 20°C to t. Resistance to oxida tion of  the samples was  investigated in a stationary atmo sphere  of  a muffle  resistance  furnace  in  the  range  of  1000 - 1800°C using crucibles made of Al2O3. The UHTC  sample was heated at a rate of 10°C/min and held at the max imum temperature for 30 min,  then the furnace was allowed  to cool by inertia to room temperature.  RESULTS AND ITS DISCUSSION  The starting powders  The main characteristics of the starting raw materials are  presented in Table 1. According to the results obtained,  for  powders of HfB2 produced by KTM,  the measured r  p is less  than the theoretical value (r  p = 6.906 vs. r t = 11.20 g/cm3).  This may indicate the presence in the powder of impurities of  boron oxide and boric acids. Together with elemental analy sis data,  it  is possible  to estimate  the  content of B2O3  in  HfB2 powders at a level of 4 to 8 wt.%. A low content of ox ygen  and  carbon  admixtures  suggests  that  the  initial  HfB2 powders were obtained by the interaction of elements  or by borothermal reduction (reduction of HfO2 by boron).  According  to XRD data, HfB2 powders  contain HfO2  (mon.) in an amount of not more than 1 wt.% (see Table 1).  No  other  crystalline  impurities  were  detected  in  the  HfB2 powders. The calculated parameters of the HfB2 crystal  lattice  of  both  powders  within  the  error  range  (a = 3.141(2) Å,  c = 3.474(2) Å)  are  in  a  good  agreement  with the literature data (a = 3.139 Å, c = 3.473 Å [1]). Elec tron microscopy showed that  the shape of  the particles  in  HfB2 powders  is  close  to  spherical with  a  diameter  of  1 - 4 mm. The particles gather in agglomerates of an indeter minate shape with dimensions of 10 - 40 mm. These data are  in good agreement with the results of laser diffraction. By the  combination  of  characteristics  (dispersity,  chemical  and  phase purity), the HfB2 powders used in this study are similar  to the powders described in [4 - 6]. The initial SiC powders  of  technical purity8 contain the permissible amount of oxy gen and iron (see Table 1).  The HP process  Intensive shrinkage under HP conditions usually began  already at 1300°C and was fully complete at 1800°C, which,  apparently,  is due to the formation of phases of reduced vis cosity in the HfO2-SiO2-B2O3  system. When reaching the  high temperature,  the actual density of  the ceramic did not  increase.  Ceramics after HP  The typical appearance of ceramics after HP, sanding and  polishing is shown in Fig. 1.  It can be seen that monolithic  ceramics were obtained successfully without visible cracks  or inclusions. According to hydrostatic weighing, the relative  density of  the ceramics was 97 - 99% (Table 2). A slightly  306  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  TABLE 1. Basic Characteristics of the Starting Powders.  Material  Provider; brand; standard  r  p, g/cm3  Phase composition**  Particle size, mm  Ssp, cm2/g (PSKh-9)  Main impurities, wt.%  HfB2  KTM;  TU 6-09-03-418-76  6.906*  HfB2, weak tr. HfO2  1 - 40, d50 = 8.86  1900 (d = 2.80)  Hf 83; B 11; Cu 0.7;  O 5.3; C 1.0  “Ekos-Ural”;  TU 6-09-03-418-76  10.914  HfB2, weak tr. HfO2  1 - 40, d50 = 9.40  —  Hf 88; B 8.0; O 3.0;  C 1.0  SiC  ZAP; KKZ, 64C, M5;  GOST 26327  —  a-SiC  2 - 18, d50 = 7.21  6675 (d = 2.78)  Cfree 0.27; Fe 0.09;  O 0.65  VAP; KKZ, 64C, M5;  GOST 26327  —  a-SiC  1 - 13, d50 = 4.97  9898 (d = 1.87)  Cfree 0.27; 0.04;  Fe 0.2; O 0.5  * May indicate a presence of up to 8 wt.% B2O3.  ** Weak tr. - weak traces (<1 wt.%).  Fig. 1. HfB2 - 30% SiC ceramics after HP and machining: left - af ter polishing; right - after sanding.  8  For HP it  is better  to use special grades of highly disperse SiC  powders.  \\x0c', 'smaller value of r  a in samples Nos. 1 and 2 can be explained  by the presence of a larger  fraction of B2O3  in the initial  HfB2 powder  (KTM). The minimal  porosity  of  less  than  0.1% was obtained for samples Nos. 1 and 4 (measured on  70 ´ 32 ´ 4.5 mm beams);  in samples Nos. 2 and 3, porosity  (~ 0.4%) was measured on 3.0 ´ 4.0 ´ 45.0 mm beams. Wa ter absorption of all samples was at the level of hundredths of  a percent.  According to the results of optical microscopy (Fig. 2),  isolated grains of SiC are present  in the ceramic matrix of  HfB2. The  average  size  of  grain  of  irregular  shape was  (4.6 ± 0.3) microns for HfB2 (KTM), (4.8 ± 0.5) mm for SiC  (ZAP),  (4.5 ± 0.3) mm  for  HfB2  (“Ekos-Ural”)  and  (2.9 ± 0.3) mm for SiC (VAP). The average grain surface area  was 13 mm2 for HfB2 (KTM) and 14.3 mm2 for SiC (ZAP),  12 mm2 for HfB2 (“Ekos-Ural”) and 5.7 mm2 for SiC (VAP).  Thus, there is no significant increase in grain size during the  selected HP mode. In sample No. 3 the formation of individ ual aggregates of SiC grains 40 - 60 mm in size was discov ered. Most likely, aggregates are formed during the drying of  the  solvent  after  sanding. Possibly,  ultrasonic  dispersion,  mixing in another  solvent and/or  the use of  surfactants, as  well as special  technologies for  removing the solvent  from  the charge (freeze-drying, the use of a rotary evaporator, etc.)  would yield a more even distribution of SiC particles in the  HfB2 matrix. The presence of a large number of SiC aggre gates in ceramics can negatively impact its mechanical prop erties [1].  Optical microscopy (up to ´1000) did not detect  any  pores, cracks or inclusions of other phases in the investigated  cuts, which indicates a good quality of the starting raw mate rial.  The  volume  fraction  of  SiC  was  estimated  as  30 - 40 vol.%, which corresponds  to the original quantity.  According to the results of XRD analysis, all samples have  identical phase composition and consist of HfB2  (hex.) and  SiC (hex. 6H with an admixture of rhomb. 57R); HfO2 is ob served in small amounts (a mixture of orthorhomb. and tetr.  accounting for no more than 1 wt.%). The presence of other  phases besides the original composition is unlikely. The frac Ultra-High-Temperature Ceramics Based on HfB2 - 30% SiC  307  TABLE 2. The Main Properties* of Sintered Ceramics (HP: 20 MPa, 1800-2000°C, exposure 30 min, duration 2 hrs).  Sample  Composition, vol.%  r  a, g/cm3  r  a/r  t, %  P, %  W, %  s fl, MPa  HV1.0, GPa  KIc, MPa·m1/2  a  20-1100, 10-6 K-1  No. 1  HfB2 (KTM) - 30SiC (ZAP)  8.40  96.0  0.09  0.01  314 - 361  17.3 ± 0.6  4.6 ± 0.5  6.6  No. 2  HfB2 (KTM) -30SiC (VAP)  8.53  96.9  0.43  0.05  281 - 702  19.6 ± 0.6  5.1 ± 0.3  —  No. 3  HfB2 (“Ekos-Ural”) - 30SiC (ZAP)  8.61  97.8  0.48  0.05  275 - 360  19.4 ± 0.2  5.9 ± 0.6  —  No. 4  HfB2 (“Ekos-Ural”) - 30SiC (VAP)  8.79  99.9  0.09  0.01  286 - 556  21.4 ± 0.9  5.5 ± 0.5  6.6  * HV1.0 at a load of 1 kg (9.8 N); KIc at a load of 3 kg at 20°C.  Fig. 2. Optical photographs (´1000) of  the surface of polished HfB2 - 30% SiC ceramics: a - d)  respectively, samples Nos. 1 - 4;  light —  HfB2, dark — SiC.  \\x0c', 'tion of SiC, based on XRD, is about 11.0 wt.%, which corre sponds to the original quantity.  TCLE (see Table 2)  is in good agreement with the data  available in the literature [3]. Within the error of measure ment  (~ 3%),  the temperature dependence of  the curves of  thermal  expansion of  samples Nos. 1 and 4 is  completely  identical.  The  average  TCLE  increases  sharply  from  5.2 ´ 10-6 to 6.6 ´ 10-6 deg-1 in the range from 20 to 200°C  and from 20 to 1100°C, respectively. This causes a reduction  in the density of  the ceramic from 8.79 to 8.61 g/cm3 at 20  and 1100°C. The values obtained are very close to the TCLE  of ceramics based on ZrB2 - 30% SiC [3, 7].  The microhardness of all samples is comparable to the  hardness of  sapphire (a-Al2O3). From Fig. 3 and Table 2,  two patterns are recognized: an increase in the relative den sity of the ceramic and a decrease in the grain size of HfB2  and SiC (at nominally the same composition)  increase the  microhardness;  the same tendency has ben previously ob served [12]. The microhardness of HfB2-SiC ceramics is sig nificantly (5 - 6 GPa) greater than that previously measured  in ZrB2-SiC UHTC [7, 8]. This can be explained by the high  chemical purity of  the  initial HfB2 powders. The obtained  value of HV ~ 20 GPa is  typical  for hot-pressed HfB2-SiC  ceramics with grain sizes of only a few microns (see Table 3  and Fig. 3)  [4 - 6, 9 - 14]. At  the same time,  in the occa sional  publications  on HfB2-SiC  ceramics  obtained  by  electropulse plasma sintering (spark plasma sintering (SPS)),  HV is  said to be  equal  to (26.0 ± 1.0) GPa  (dHfB2 ~ 2 mm,  dSiC ~ 1 mm) in [13] and even (27.0 ± 0.6) GPa ( dHfB2 ~ 6.2 mm,  dSiC ~ 1 mm) in [9], however, these HV values may be signifi cantly overestimated due to the significant  thermal stresses  characteristic of  the SPS hardened ceramics.  It  should be  noted that the grain size in ceramics produced for the present  study  is  about  2 times  larger  than  reported  in  Table  3  [4 - 6, 9 - 14]. Thus,  the  ceramics produced in this work  may have a significant reserve for further increase in micro hardness.  In  all  ceramic  samples, KIc > 4.6 MPa·m1/2, which  is  comparable  to the  crack resistance of  ceramics based on  ZrB2-SiC sintered by HP and SPS [7, 8]. At  the same time,  our KIc data significantly exceed a number of values encoun tered in the literature on HfB2-SiC ceramics (see Table 3).  Further increase in microhardness and KIc is possible with a  significant  reduction in the grain size of HfB2 and SiC to  1 mm (or less) or by adding an additive, for example B4C [5]  (which may worsen oxidation resistance and reduce thermal  conductivity).  Fractographic analysis  The key structural characteristic of composite ceramics  based on HfB2-SiC is the static limit of flexural strength s  fl.  In  the  samples  after  mechanical  treatment,  s  fl  was  300 - 500 MPa  at 20°C (see Table 2). The upper  limit of  strength of  samples Nos. 2 and 4 with fine SiC (VAP)  is  higher  than for  samples Nos. 1 and 3 with relatively large  SiC (ZAP). Samples usually broke into 2 - 3 pieces. At  the  same time, strength did not appear  to depend on the aniso tropy of samples (along and perpendicular to the pressure ap plication axis during the HP). A typical value of s  fl  for  the  hot pressed UHTC HfB2 - SiC is 500 - 700 MPa (see Ta ble  3)  [5 - 6, 9 - 14]. However,  in  [11], s  fl  values  up  to  1 GPa have been reported.  According to the results of  fractographic analysis  (Ta ble 4), the reason for the large spread of strength values is the  presence  of  large  defects  containing molybdenum up  to  150 mm in size in the near-surface volume of  the ceramic.  The defect  is usually at a distance of about 20 mm from the  surface. At  the same time, according to EDXMA,  sodium  and potassium are often found in areas around the defects.  Another  type of defect  is cracks starting at  the surface and  penetrating into the  sample  to a depth of 20 - 100 mm.  It  should be noted that no copper or zinc (the main components  308  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  TABLE 3. Comparison of the Characteristics of Hot-Pressed Ceramics According to the Authors and Literature Data.  Composition, vol.%  d  HfB  2  , mm  dSiC, mm  r  rel, %  s  fl, MPa  HV, GPa  KIc, MPa·m1/2  Reference  HfB2 - 30SiC  4.5  2.9  99.9  300 - 700  21.4 ± 0.9  5.5 ± 0.5  According to the authors of [4]  HfB2 - 26SiC  5.0  3.0  99.8  692 ± 58  18.3 ± 0.3  5.23 ± 0.1  [12]  HfB2 - 20SiC  2.5  2.0  98.6  526 ± 86  19.5 ± 0.8  3.95 ± 0.4  [6]  HfB2 - 20SiC  2.0  ~1  —  993 ± 90  20.2 ± 0.1  6.29 ± 0.65  [11]  HfB2 - 20SiC - 2B4C  3.5  ~2  99.6  771 ± 30  21.6 ± 0.8  7.06 ± 0.4  [5]  HfB2 - 20SiC - 5WC  1.5  1.5  99  544 ± 135  22.3 ± 1.5  3.76 ± 0.7  [6]  HfB2 - 19SiC - 6Si3N4  ~2  ~1  95.5  560 ± 100  20.4 ± 0.6  4.3 ± 0.1  [14]  HfB2 - 20SiC - 2TaSi2  ~2  ~1  —  665 ± 75  —  3.6 ± 0.5  [10]  HfB2 - 30SiC (SPS)  ~2  ~1  >99  590 ± 50  26.0 ± 1.0  3.9 ± 0.1  [13]  HfB2 - 20SiC (SPS)  6.2  ~1  >99  620 ± 50  27.0 ± 0.6  5.0 ± 0.4  [9]  \\x0c', 'of brass - rubbing sieve screen material) were detected in the  defects of the investigated ceramics.  According to T.  I. Serebryakova [15],  in zirconium and  hafnium diborides obtained by borothermal reduction, Al, Ni  and Fe (up to 0.28 wt.%), C (up to 0.25 wt.%), Ti  (up to  0.13 wt.%), Mg (up to 0.15 wt.%) and S (up to 0.03 wt.%)  are the typical admixtures. According to G. G. Gnesin [16],  technical grade  silicon carbide may contain impurities of  complex composition in the form of microscopic inclusions.  Most  inclusions are silicides and carbides of iron, as well as  more complex systems (Si-Fe-Mn-Ni, Al-Si-Ca, etc.). The  gross content of Fe, Al  is at  the level of tenths of a percent,  and of Mg, Ti, Ca, Zr, Ni, Cr - at the level of hundredths of a  percent. Metallic impurities are introduced into SiC from the  raw material:  silica, carbon ash and common salt  [16], as  well  as  from grinding bodies used for  crushing powders.  Thus, the admixture of molybdenum is atypical for the com ponents of UHTC. Perhaps  this  is an accidental  (uninten tional) impurity. It should be noted that there are a number of  works in which the elemental molybdenum [17] and its com pounds (for example, MoSi2) [1] are purposefully introduced  into the UHTC to improve its properties (strength, hardness,  KIc, etc.). In addition, it is recommended to use molybdenum based alloys as special solders for UHTC compounds [18].  In the preliminary oxidation of  the  samples  and their  subsequent  testing  at  room temperature  (Table 4), s  fl  in creased (338 - 842 MPa). The limit of  flexural  strength of  sample No. 4 at 20°C after preliminary oxidation at 1400°C  (60 min) and at other temperatures (30 min) is given below:  T,°C .  .  .  .  .  .  .  .  . Without thermal  treatment  1400  1500  1600  1700  1800  s  fl, MPa .  .  .  .  .  .  286 - 556  603 - 790  690  459  206  235  In this case,  the defects initiating the destruction are ex clusively in the bulk of  the ceramic (see Table 4). The test  Ultra-High-Temperature Ceramics Based on HfB2 - 30% SiC  309  Fig. 3. Vickers microhardness of samples Nos. 1 - 4: hollow sym bols — data of  the authors of  the article (one symbol — 3 tests);  solid symbols — literature data. For convenience, data points are  connected by lines.  TABLE 4. Results of Fractographic Analysis of the Key Samples of the HfB2-SiC UHTC*.  Sample  s  fl, MPa  Type of defect initiating breaking  Location of the defect  Distance to chamfer, mm  Defect size, mm  Note  No. 1  329  Mo inclusion  At the surface  Near the chamfer  70 ´ 70  Without oxidation  No. 2  281  Mo, Fe, O inclusion  At the surface  1500  50 ´ 60  Without oxidation  No. 2  500  Zr, Ca, O inclusion  At the surface  850  30 ´ 40  Without oxidation  No. 4  275  Mo, Fe, S, O, Cl inclusion  At the surface  1300  90 ´ 90  Without oxidation  No. 4  499  Al, O inclusion  At the surface  1100  30 ´ 40  Without oxidation  No. 4  302  Mo, Fe, Cr, O, C inclusion  In the bulk, 250 mm  from the face  1700  150 ´ 200  Without oxidation  No. 1  342  Mo, S inclusion  In the bulk, 400 mm  from the face  1200  120 ´ 150  With oxidation  No. 2  624  Mo, S inclusion  In the bulk, 450 mm  from the face  1400  50 ´ 70  With oxidation  No. 3  338  Area with the presence of S  In the bulk, 360 mm  from the face  1700  90 ´ 100  With oxidation  No. 4  697  Zr inclusion  In the bulk, 270 mm  from the face  1200  50 ´ 60  With oxidation  No. 4  790  Zr inclusion  In the bulk, 290 mm  from the face  1200  40 ´ 40  With oxidation  * The composition of impurities according to EDXMA. The characteristic lines of Hf and Si are similar. The determination of light elements  (from boron to oxygen) is difficult.  \\x0c', 'samples in this case broke into 5 - 9 pieces. The positive ef fect of heat treatment was repeatedly noted earlier for a num ber  of  refractory  systems: TiN-TiB2  [19], ZrB2-SiC and  HfB2-SiC [9, 20]. During calcination in air, a continuous ox ide film forms, which heals microcracks on the surface of the  ceramic.  In  addition,  heat  treatment  partially  removes  microstresses. Thus,  the intrinsic strength of ceramics (with out  large random defects and defects induced during sand ing) is quite high.  It can be assumed that improving the chemical and phase  purity of the original HfB2 and SiC powders, further reduc ing the SiC grain size (possibly, by switching from M3 to  M1 grade) and optimization of the batch preparation process  and HP parameters will increase the strength of the ceramics.  To get rid of the largest  inclusions (defects) of ceramics,  the  following steps can be recommended: to conduct preliminary  sieving of silicon carbide;  to use tungsten carbide balls with  dimensions less than 5 mm for milling;  to use a sieve with a  smaller cell size for sieving after grinding;  to lower  the HP  temperature to 1800 - 1900°C and increase exposure at  this  temperature.  Resistance to oxidation  Curing in air at up to1000°C (inclusive) does not change  the appearance of  the samples and their mass.  In the range  from 1100 to 1300°C a dark matte coating forms on the sam ples. No  interactions were  observed  (adhesion)  between  HfB2 - 30% SiC samples and Al2O3and SiC-based ceram ics during joint  firing in air  in the temperature range up to  1300°C inclusive.  When exposed to air  for 1 hour  in a muffle furnace at  1400°C,  the weight gain of  the UHTC was (0.099 ± 0.007)  wt.% (0.75 ± 0.09 mg/cm2);9 according to electron micros copy,  the  thickness of  the  resulting uniform glassy oxide  layer is 3 - 4 mm.10 The resulting film acts as a barrier to fur ther oxidation; with repeated oxidation under the same con ditions,  there was no change in mass. The thickness of  the  “depleted SiC”  layer was  not more  than  15 mm. Below,  20 mm deep,  is unoxidized ceramics (except  for  the case of  cracks on the surface). At 1500 - 1600°C (30 min exposure),  single inhomogeneities and swelling of the oxide film are ob served on the  samples  along with an increase  in mass of  0.18 - 0.26%, or 1.5 - 1.9 mg/cm2. At 1500°C,  a  layer of  hafnium oxide (Fig. 4a ) with a thickness of 1 - 2 mm already  appears beneath the glassy layer. At 1600°C, the thickness of  the  glassy  layer  is  5 - 6 mm,  and  of  the HfO2  layer  is  6 - 7 mm (Fig.  4b ).  At  1700°C, mass  gain  is  0.37%  (~2.6 mg/cm2), and the thickness of the HfO2 layer increases  to 25 - 27 mm (Fig. 4c ). At 1800°C,  the mass gain is 0.44%  (~3.2 mg/cm2),  the oxide layer acquires a complex structure  with numerous bubbles, microcracks, detachments and inclu sions of hafnium oxide. The total thickness of the loose oxide  layer reaches 30 mm (Fig. 4d ). Grains of hafnium oxide up to  310  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  Fig. 4. Electron microscopy of the oxide layer and the fracture of HfB2 - 30% SiC ceramics (sample No. 4) after oxidation at different tempera tures for 30 min in air at 1500 (a), 1600 (b), 1700 (c) and 1800°C (d).  9  Statistics for 12 samples measuring 3.0 ´ 4.0 ´ 45.0 mm. Despite  the different densities, no difference in mass gain was observed  between the UHTC samples in terms of resistance to oxidation.  10  The calculated density of such a coating is ~ 2.0 g/cm  3  , as in the  case  of  glass with  a  composition  of  65% SiO2 - 35% B2O3.  For  reference:  the density of amorphous SiO2 and B2O3 is 2.21  and 1.87 g/cm  3  , respectively.  \\x0c', '5 - 7 mm in size are oriented in columns. A glassy melt  is  found in the cavities between the columns. The thickness of  the “depleted SiC” layer (between the unoxidized bulk of the  ceramic and the HfO2  layer)  is difficult  to estimate.  In any  case,  its thickness is no more than 20 mm at 1500 - 1600°C  and  30 - 40 mm  at  1700 - 1800°C.  Oxidized  samples  (at ³1600°C) acquire a lighter gray color with an increase in  temperature.  The  observed  picture  at  the  qualitative  level  and  the  quantitative  values  obtained  are  in  good  agreement with  those previously reported in the literature. For example,  in  the dissertation [21],  the mass gain is 0.14% in HfB2 - 30%  SiC after aging at 1400°C, in [9] for HfB2 - 20% SiC ceram ics after oxidation at 1400°C for 1 hour  it  is  (0.60 ± 0.03)  mg/cm2,  and  in  [10],  after  oxidation  at  1450°C,  it  is  0.79 mg/cm2. In [12], after oxidation of HfB2 - 26% SiC ce ramics (r  a = 9.12 g/cm3) at 1500°C for 30 min, the thickness  of the oxide film is about 3 mm (thickness of the “lean SiC”  layer  is ~ 10 mm). When tested at 1800°C for 5 hours  in a  muffle furnace,  the increase in the mass of  the HfB2 - SiC  ceramic (r  f ~ 7.2 g/cm3, r  rel ~ 90%) can be 0.9% [22]. Heat  resistance can be improved by adding TaSi2 [10], MoSi2 [1]  and La2O3 additives to the HfB2 - SiC based ceramic [9, 20].  Each additive is optimal  for a particular  temperature range  and operating mode.  CONCLUSION  Using entirely domestic raw materials,  it  is possible to  obtain a dense (up to 99.9%) ceramic of  the composition  HfB2 - 30 vol.% SiC by hot pressing. This ceramic has high  hardness  (up to 20 GPa) and KIc  (up to 5.9 MPa·m1/2),  its  flexural strength ranges from 300 to 700 MPa at 20°C.  The  authors  thank N. A. Golubeva, G. M. Bagreeva,  P. Yu. Yakushkia, and A. P. Metleva for their help and assis tance in carrying out the research work.  REFERENCES  1. Ultra-High Temperature Ceramics. Materials for Extreme Envi ronment Applications, Ed. by W. G. Fahrenholtz, E. J. Wuchina,  W. E. Lee, and Y. Zhou, Wiley, New Jersey (2014) 441 p.  2. T. H. Squire and J. Marschall, “Material property requirements  for analysis and design of UHTC components in hypersonic ap plication,” J. Europ. Ceram. Soc., 30, 2239 - 2251 (2010).  3. M. Mallik, A. J. Kailath, K. K. Ray, and R. 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Vaganova, “Features of sintering of refractory ceramics  based on HfB2 by hybrid spark plasma sintering” [in Russian],  Zh. Prikl. Khim., 88(3), 379 - 386 (2015).  Ultra-High-Temperature Ceramics Based on HfB2 - 30% SiC  311  \\x0c']"
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  "_id": 279,
  "PDF": "Ultra-high-temperature testing of sintered ZrB2-based ceramic composites in atmospheric re-entry environment.pdf",
  "Text": "['International Journal of Heat and Mass Transfer 156 (2020) 119910   Contents lists available at ScienceDirect   International Journal of Heat and Mass Transfer   journal homepage: www.elsevier.com/locate/hmt   Ultra-high-temperature testing of sintered ZrB 2 -based ceramic   composites in atmospheric re-entry environment   Stefano Mungiguerra  , Giuseppe D. Di Martino  a , Anselmo Cecere  a , Raffaele Savino  b , Laura Silvestroni  b , Diletta Sciti  Luca Zoli   a ,   a ,   ∗  b   a University of Naples “Federico II”, Department of Industrial Engineering, P.le Vincenzo Tecchio 80, 80125 Napoli, Italy   b National Research Council, Institute of Science and Technology for Ceramics, Via Granarolo 64, 48018 Faenza, Italy   a r t i c l e   i n f o   a b s t r a c t   Article history:   Received 24 December 2019   Revised 29 April 2020   Accepted 2 May 2020   Available online 18 May 2020   Keywords:   An experimental campaign has been carried out to characterize a new class of Ultra-High-Temperature   Ceramic Matrix Composites for near-zero ablation Thermal Protection Systems. Small-sized specimens,   with ZrB 2 -based matrix and different carbon fiber architectures, were exposed to a simulated air super  sonic flow generated by an arc-jet wind tunnel, achieving specific total enthalpies up to 20 MJ/kg and cold   wall fully catalytic heat fluxes over 5 MW/m  2 , in an aero-thermo-chemical environment representative of   atmospheric re-entry. Ablation rates were estimated by means of mass and thickness measurements be  Ultra-high-temperature ceramic matrix   fore and after testing, demonstrating an excellent performance of the developed materials. Surface tem  composites   Arc-jet wind tunnel testing   Near-zero ablation   Computational fluid dynamic simulation   Temperature jump   peratures were monitored by means of infrared pyrometers and a thermo-camera, and during all the   tests a spontaneous temperature jump was observed, with temperatures that reached values over 2800   K at the steady state. Post-test microstructural analyses revealed the formation of a porous oxide layer   with a thickness of few hundred microns, mainly consisting of zirconia, with substantial removal of both   SiC and carbon fibers. Below the oxide, the bulk material was unaffected. Com putational Fluid Dynamics   simulations allowed rebuilding the thermo-fluid-dynamic and chemical flow field. Moreover, it was pos  sible to propose an innovative correlation of the temperature jump with an increased catalytic activity   and a dramatic reduction of the thermal conductivity of the oxide layers forming on the exposed part   of the sample, which anyway had a key role in preserving the unoxidized bulk materials at reasonable   temperatures.   © 2020 Elsevier Ltd. All rights reserved.   1. Introduction   over 20 0 0 °C [3-5] . Nevertheless, some issues related to poor   oxidative behavior and mechanical properties (damage tolerance,   New-generation hypersonic and reusable re-entry vehicles set   fracture toughness, thermal shock resistance) of single and multi  increasingly demanding requirements for the development of high  phase UHTCs at high temperatures limit the applicability of these   performance Thermal Protection Systems (TPS), due to the chal  materials. The introduction of SiC or other silicon based ceramics   lenges of extremely harsh aero-thermo-dynamic conditions charac  as minority phase, in the form of particles, short/long fibers or   teristic of atmospheric re-entry, including hypersonic Mach num  whiskers [6-10] , in the main refractory ceramic has been often   bers, temperatures above 20 0 0 °C, the activation of gas dissocia  proposed to improve damage tolerance and oxidation resistance at   tion/recombination reactions at extremely low oxygen partial pres  intermediate temperature, essentially thanks to the formation of a   sures, which can substantially enhance the heat flux on the ex  low-viscosity borosilicate glass protective scale [11-13] . The most   posed surface of the spacecraft [ 1 , 2 ].   recent frontiers in a research oriented to high Technology Readi  Over   the   last   decades,   research   identified   Ultra-High  ness Level (TRL) applications of the UHTC technology to aerospace   Temperature Ceramic   (UHTC) materials, based on   transition   involve the enhancement of mechanical properties by introducing   metals carbides and diborides, as potentially promising candidates   short and continuous Carbon Fiber reinforcements in a UHTC   for these applications, especially in light of their high melting   matrix, leading to the definition of the Ultra-High-Temperature   temperatures, strength and ablation resistance at temperatures   Ceramic Matrix Composites (UHTCMCs) [ 14 , 15 ]. Recently, flexural   ∗  Corresponding author.   E-mail address: stefano.mungiguerra@unina.it (S. Mungiguerra).   https://doi.org/10.1016/j.ijheatmasstransfer.2020.119910   0017-9310/© 2020 Elsevier Ltd. All rights reserved.   strength values as high as 450 MPa and 20 0-30 0 MPa were   collected at 1500 °C and 2100 °C,   respectively, demonstrating   the excellent performance of UHTCMCs based on Carbon fibers   \\x0c', '2   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   preforms and HfC, ZrC, TaC and ZrB 2   matrices [16-18] . All the high   temperature values were always higher than the room tempera  ture values. Moreover, UHTCMCs based on carbon preforms and   ZrB 2   matrix displayed also excellent thermal shock resistance [18] ,   with 85% of the pristine flexural strength retained after water   quenching from 1500 °C to room temperature.   The overall objective   is developing   large ultra-refractory   aerospace transportation systems components with outstanding   ablation resistance and enhanced mechanical properties and reli  ability [19-22] . To achieve this goal, testing in a relevant environ  ment is required to properly characterize the ceramic materials in   conditions representative of the real flight applications. For atmo  spheric re-entry TPS, the most suitable facilities are supersonic and   hypersonic arc-jet wind tunnels [ 23 , 24 ].   Within this framework, University of Naples “Federico II” (UN  INA) and the Institute of Science and Technology for Ceramics of   the Italian National Research Council (CNR-ISTEC) are involved in   the Horizon 2020 European C  3 HARME research project, focused on   the development of a new class of UHTCMCs for near-zero ablation   Thermal Protection Systems [25] . An extensive experimental char  acterization campaign has been carried out on in the UNINA arc  jet wind tunnel, where atmospheric re-entry conditions are repro  duced at maximum flow total enthalpies higher than 20 MJ/kg, su  personic Mach number and temperatures over 20 0 0 °C in a gas at  mosphere with high concentration of atomic oxygen. Non-intrusive   diagnostic equipment, including two-color pyrometers and an in  frared thermo-camera, is employed to monitor the surface temper  ature of the samples. The ablation rates of the samples after the   exposure to the aero-thermo-chemically aggressive flow are also   estimated by mass and thickness measurements.   We previously published the results of an arc-jet wind tunnel   characterization campaign carried out on a batch of two UHTCMC   samples with ZrB 2  -10 vol% SiC as matrix and ~50 vol% of continu  ous coated PAN-based carbon fibers as reinforcing phase [26] . The   present work extends the results to 3 more samples, all based on a   UHTC matrix with ZrB 2  as major component and SiC and Y 2   O 3   par  ticles as minority phase, reinforced with uncoated continuous or   chopped ultra-high modulus pitch-based carbon fibers. Compared   to the previous work [26] , in this one, we notably improved the   manufacturing of UHTCMCs. The matrix was nearly pore-free, the   matrix phase distribution in the fiber preform was more homoge  nous. These two features are of fundamental importance to guar  antee a better protection of the fibers by the UHTC matrix during   oxidation and hence arc-jet tests, as explained in [15] .   The outcomes of the experimental activities are presented and   discussed, also in light of the post-test characterizations carried   out to investigate the features of the materials microstructures af  ter the exposure to the atmospheric re-entry environment. More  over, the experimental results are complemented by Computational   Fluid Dynamics (CFD) simulations, employed to allow accurate pre  diction not only of the thermo-fluid-dynamic flow field around the   test articles, but also of the thermal behavior of the materials sam  ples, including an investigation of the effect of material properties,   such as thermal conductivity and catalycity.   Specifically, a rapid temperature increment during the highest  enthalpy steps, phenomenon known as “temperature jump” in the   relevant literature regarding UHTCs and SiC-based ceramics, has   been observed [27-29] . This paper intends to propose a thorough   and detailed analysis of the materials aero-thermo-dynamic behav  ior at ultra-high temperatures in a representative re-entry envi  ronment, aiming to provide a comprehensive interpretation of the   temperature jump, correlating the outcomes of infrared tempera  ture measurements, post-test microstructural analyses and numeri  cal simulations to highlight the parameters which mainly affect the   heat transfer from the flow to the ceramic. A deep understanding   of the materials response is indeed an important step in the path   Table 1   List of matrix compositions and fiber volume fraction (FVC) of the in  vestigated UHTCMCs.   Sample ID   UHTC, vol%   FVC, vol%   Fiber type   ZS5Y-SF   ZrB 2   + + +   5SiC   + + +   5 Y 2 O 3   40   short, XN60   ZS10Y-SF   ZrB 2    10SiC    5 Y 2 O 3   45   short, XN60   ZS10Y-LF   ZrB 2    10SiC    5 Y 2 O 3   35   long, XN80   towards a reliable engineering application of these novel TPS tech  nologies.   2. Materials and methods   2.1. Material processing   Different matrix formulations and carbon fibers architectures   have been compared. UHTC matrices containing different amount   and type of additives (SiC and Y 2   O 3   ) were prepared by conven  tional wet powder milling technique, Table 1 . SiC is a common   secondary phase added into UHTC matrices to increase the oxi  dation resistance and mechanical properties, Y 2   O 3   is introduced   with the aim of improving the oxide stability, i.e. preventing the   ZrO 2  martensitic transformation and corresponding volume expan  sion upon cooling. Regarding fiber, both random chopped fibers   and continuous fibers alternately arranged with 0 ° and 90 ° orien  tation have been considered.   Short fibers were added to the UHTC powders and mixed for   24 hours. The slurries were then dried in a rotary evaporator and   de-agglomerated. Long fiber-reinforced materials were obtained by   aqueous slurry infiltration, stacking of the Cf foils with a 0 °/90 °  configuration and drying in air.   Composites were densified to full density after sintering at   1900 °C as described in [ 30 , 19 ]. Typical images of the microstruc  tures taken by scanning electron microscopy (FE-SEM, Carl Zeiss   Sigma NTS Gmbh, Oberkochen, DE) are displayed in Fig. 1 . All   composites had large defect-free microstructure. Short fibers, with   length around 150 μm, were homogeneously dispersed into the   ZrB 2  -based matrix with preferential orientation along the xy axis,   Fig. 1 (a). Long-fiber reinforced materials presented some perpen  dicular micro-cracks between the fabric layers, but good fiber infil  tration and 0 ° to 90 ° layers adhesion was achieved.   2.2. Arc-jet facility   Experimental tests were carried out in the arc-jet wind tun  nel available at the University of Naples “Federico II”, named SPES   (Small Planetary Entry Simulator). This is an open circuit, continu  ous wind tunnel where a nitrogen plasma can be generated by an   industrial torch able to operate at powers up to 80 kW and mixed   to a secondary cold oxygen flow in order to simulate earth atmo  spheric composition. A converging-diverging nozzle is employed to   expand the hot mixture to a nominal supersonic Mach number   equal to 3. A detailed facility description can be found in [31] .   Small-sized material samples can be placed downstream the   nozzle exit section by a dedicated thermally protected supporting   mechanism, in order to expose them to the supersonic dissociated   air flow, and characterize materials response to the extreme aero  thermo-chemical environment [32] . The nominal samples design   for the present test campaign is that typically used in the arc-jet   facility [28] and is displayed in Fig. 2 . The UHTCMC samples were   placed at a distance of 1 cm from nozzle exit.   In the present test campaign, samples were exposed to a su  personic flow generated by the expansion of a high enthalpy gas   mixture of nitrogen (0.8 g/s) and oxygen (0.2 g/s). During the test,   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   3   Fig. 1. Sample pictures (top right frame) and typical microstructure of as-sintered UHTC reinforced with a) short fiber or b) long fiber with 0 °/90 ° architecture.   Fig. 2. Nominal design of UHTCMC samples. Dimensions in mm.   Table 2   Specific total enthalpies at different steps during the tests.   Step   1   2   3   4   5   6   7   8   H 0 [MJ/kg]   7.0   8.5   10   12   14   16   18   20   the arc power of the plasma torch and consequently the total en  thalpy of the flow were gradually increased through successive in  crements, leading correspondingly to an increase of pressure and   temperature. A quasi-stationary condition generally occurs when   the maximum value of temperature is reached during the last steps   of the test. In the tests discussed in the present work, the nominal   duration of each step was 30 seconds, except for the last step of   120 seconds. The specific total enthalpy is obtained, with an un  certainty around 10%, through an energy balance at the exit of the   plasma torch, based on measurement of temperature and flow rate   of cooling water [31] . The values of specific total enthalpy corre  sponding to each power step are reported in Table 2 . After reaching   the maximum enthalpy level, the arc power is gradually decreased   until facility shutdown.   The surface temperature of the samples was continuously mea ±1% instrumental accuracy) by digital two-color pyrometers  sured (   (Infratherm ISQ5 and IGAR6, Impac Electronic GmbH, Germany) at   an acquisition rate of 100 Hz. In addition, an infrared (IR) thermo  camera (TC, Pyroview 512N, DIAS Infrared GmbH, Germany) allows   for the evaluation of the temperature distribution over the sam  ple surface. The ISQ5 pyrometer exploits two overlapping infrared  μm and 0.97-1.15  μm to measure   wavelength bands at 0.7-1.15   the actual temperature from 1273 K up to 3273 K. The IGAR6 py  rometer operates in the bands 1.5-1.6 μm and 2.0-2.5 μm to return   the sample temperature in the range 523-2273 K. The measure  ment area of the ISQ5 pyrometer is approximately a round spot   3.3 mm in diameter. The thermo-camera is able to detect temper  ature in the range 873-3273 K and it operates in the spectral range   from 0.8 to 1.1 μm. The temperature measurement of the IR-TC is  ελ in the instrument operat  dependent on the surface emissivity   ing spectral range. In order to set the correct value of the spectral   emittance, the TC measurement is compared to the actual temper  ature detected by the pyrometers in the two-color mode, and the   value of   ελ  is adjusted until the TC gives back the same temper  ature as the pyrometers. High-Definition videos of the tests were   recorded by means of a Camera Flea3 1.3 MP Color USB3 Vision  × 1048 and a frame rate equal to 25 fps.  with a resolution of 1328   A balance (1 mg accuracy) and a precision caliper (0.01 mm   accuracy) were used to measure mass and thickness losses of the   specimen. In the post-processing analyses, two erosion rates were   calculated: one based on the loss of mass (assuming uniform con  sumption of the sample in the axial direction) and the other eval  uated by the thickness measurement made by the caliper.   2.3. Numerical models   In order to properly characterize the flow field inside the facil  ity and to rebuild the thermal behavior of the samples, Computa  tional Fluid Dynamics (CFD) simulations were performed. In partic  ular, steady-state simulations of the flow field are performed, solv  ing the Reynolds-Averaged Navier-Stokes equations for a turbulent   multi-reacting gas mixture with five species (i.e. O, O 2  , NO, N and   N 2   ), in chemical non-equilibrium, considering finite rate chemical   reactions, with the reaction rate constants specified by the Arrhe  nius law [33] .   The flow simulations were performed in two steps. First the 2D  axisymmetric computational domain shown in Fig. 3 (a) was con  sidered, including the mixing chamber, where N 2   coming from the   torch and cold O 2   are mixed, and the converging-diverging nozzle,   where the high-enthalpy simulated air flow is expanded to super  sonic conditions. The total temperature and chemical composition   of the gas coming from the torch, at nitrogen inlet, were evaluated   by means of the NASA CEA (Chemical Equilibrium with Applica  tions) software [34] in order to match the total specific enthalpy   corresponding to the desired value of the torch arc power. A ra  dial mass flow inlet provides cold oxygen injection. Nozzle water  cooling was taken into account by setting a temperature bound  ary condition on the nozzle walls (T   =   400 K). The main output   of these simulations were the thermo-fluid-dynamic and chemical   conditions achieved at nozzle exit section, which was modelled as   a supersonic pressure outlet and corresponds to the inlet of the   test chamber.   In the second simulation step, the aero-thermo-dynamic flow   field around the sample and its supporting system was calculated,   using the computational grid shown in Fig. 3 (b). The main quanti  ties of interest (temperature, pressure, velocity and chemical com  position) calculated at the nozzle exit section in the previous step,   were assigned to a boundary representing the inlet of the flow into   the test chamber (bottom left of Fig. 3 (b)). On the top and rear   boundaries of the test chamber domain, a pressure-outlet condi  tion corresponding to the experimental vacuum environment was   set.   \\x0c', '4   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   Fig. 3. Mesh for the CFD simulations: a) mixing chamber and nozzle; b) test chamber, including test sample and solid supporting system.   =  The surfaces of the solid components were first treated as cold   The thermal histories of the different sam ples, measured by the   walls (T    300 K), for the calculation of the convective heat flux.   ISQ5 pyrometer, are plotted in Fig. 4 .   Subsequently, the thermal behavior of the UHTCMC sample was   The stepwise increase in temperature is associated to the power   simulated by solving the energy equation inside the solid compo  increase procedure. At the maximum enthalpy level, the tempera  nents, considering thermal coupling between fluid and solid do  tures of all the samples reached values close to 2700 K and even   main, by means of temperature and heat flux continuities at the   overpassed 2800 K for the only sample tested at the highest power   interfaces.   condition, ZS10Y-SF. In particular, it was clear that a sudden rise   Due to the high concentration of dissociated species in the flow,   in temperature (herein defined as “temperature jump”) of several   particular focus was given to the effect of surface catalycity, which   hundred degrees occurred during the highest-enthalpy steps, even   enhances the heat transfer to the sample. The non-catalytic (NC)   at a constant arc power, after a steady state condition had appar  condition was specified by assigning, for each species, a zero diffu  ently been reached. The jump happened for all the samples when   sive flux in the direction normal to the specimen surface. The fully   catalytic condition (FC) was set by assigning, on the sample sur  the flow specific total enthalpy was 18 MJ/kg. For sample ZS5Y-SF   the jump occurred around t    150 s, few seconds after the surface   =  face, the species mass fractions corresponding to a complete re  temperature reached a stable value over 2250 K, while for samples   =  =  combination of atoms into molecules. Intermediate values of the  catalytic recombination coefficients  γ w [ 35 , 36 ] can also be consid  ered, according to the model described in [26] . In this paper, to   ZS10Y-SF and ZS10Y-LF, the momentary equilibrium temperature   was around 2150 K (achieved a t    160 s and 180 s, respectively).   For sample ZS10Y-SF the torch arc power was further increased to   =  reduce the number of unknown quantities, it has been assumed   the maximum value (H 0    20 MJ/kg) at t    180 s, when the jump   that the catalytic recombination coefficients are the same for the   had already been triggered and the surface temperature was 2270   two atomic species, nitrogen and oxygen.   K, resulting in an increased slope of the temperature time pro  A more detailed description of the numerical models employed   file and a surface temperature exceeding 2800 K at the end of the   in this work can be found in [ 26 , 37 ].   3. Experimental results   3.1. Arc-jet testing   Table 3 summarizes the test conditions to which each sample   was subjected, in terms of the enthalpy steps described in Table 2 .   Table 3   Test conditions achieved for each sample.   Sample ID   Enthalpy steps   T max achieved, K   ZS5Y-SF   ZS10Y-SF   ZS10Y-LF   2-7   1-8   2-7   2690   2850   2700   heating phase. The temperature jump phenomenon is typically ob  served for UHTCand SiC-based composites [27-29] and can be as  sociated to a drastic change in surface chemistry, as will be widely   discussed later on.   After the tests, all samples heads appeared completely oxidized,   with a white layer, mainly composed of ZrO 2  , as discussed below,   covering the surface. For samples described earlier [26] , this layer   had been found to be porous, fragile and with tendency to spall   off. In the present test campaign, the samples preserved instead   the original shape and a perfect structural integrity, despite the   clearly noticeable signs of surface oxidation (see Fig. 5 ). The long  fiber layered architecture of samples ZS10Y-LF was still observable   (see Fig. 5 (c)).   The average massand thickness-based erosion rates are re −4 mm/s.  ported in Table 4 and all of them are on the order of 10   Balance between oxygen inclusion and C, Si and B volatilization   upon sample oxidation resulted in a net, although limited, mass   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   5   Fig. 4. Temperature histories of the samples, measured by the ISQ5 pyrometer.   Fig. 5. Pictures of samples after test: a) ZS5Y-SF, b) ZS10Y-SF, c) ZS10Y-LF.   Table 4   Mass and thickness data before and after the test.   Initial mass   Final mass   Average erosion rate (mass)   Initial thickness   Final thickness   Average erosion rate (thickness)   ZS5Y-SF   7.360 g   ZS10Y-SF   ZS10Y-LF   6.709 g   7.374 g   7.270 g  −4 mm/s  3.0  \\x1610   6.553 g  −4 mm/s  5.6  \\x1610   7.300 g  −4 mm/s  2.5  \\x1610   4.90 mm   5.01 mm   4.98 mm   5.00 mm  −4 mm/s  -3.2  \\x1610   5.05 mm  −4 mm/s  -1.3  \\x1610   5.01 mm  −4 mm/s  -1.0  \\x1610   loss. It is interesting to observe that, on the other hand, oxidation   uid glassy oxide phase from the front surface along the side of   led to a thickening of samples heads, resulting in a negative abla  the sample by the shear stresses induced by the supersonic flow.   tion rate based on thickness measurement. Since no other samples   A qualitative visualization of the phenomenon is possible, watch  dimensions had significant variations with respect to the nomi  ing the IR video provided in the Supplementary material. Evidence   nal values ( Fig. 2 ), only thickness data are here reported. Based on   of the waves-of-radiance can be observed also in Fig. 6 and Fig. 7 ,   these measurements, it can be concluded that all the samples ex  which are referred, as examples, to samples ZS5Y-SF and ZS10Y  perienced a slight volume increase despite the net mass loss.   LF, respectively. The diagrams on the left ( Fig. 6 (a) and 7(a) ) show   The IR video of test on sample ZS5Y-SF showed that, few sec  the temperature axial profiles measured by the TC on the side sur  onds before the temperature jump, an unsteady evolution of the   face of the samples, each curve representing a specific time instant   irradiated power appeared on the side surface, resulting in a wavy   at the earliest stages of the temperature jump. The diagrams on   oscillation of the surface temperature. A comparable phenomenon   the right ( Fig. 6 (b) and 7(b) ) report instead the time evolution of   has been already reported by Monteverde et al. [31] , who defined   the temperature at different locations along the surface. The curves   it as waves-of-radiance and correlated it to the transport of a liq  have been obtained assuming a constant spectral emissivity along   \\x0c', '6   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   Fig. 6. ZS5Y-SF: a) Temperature axial profiles, measured by IR-TC before temperature jump; b) Time evolution of the IR-measured temperature at different points along the   sample side surface.   Fig. 7. ZS10Y-LF: a) Temperature axial profiles, measured by IR-TC before temperature jump; b) Time evolution of the IR-measured temperature at different points along the   sample side surface.   the whole surface, so they should be more correctly interpreted   as surface radiance profiles, rather than actual temperature dia  grams. Two features can be remarked. First, the axial profiles of   Figs. 6 and 7 (a) are not straight, but temperature oscillations can   be noticed, up to around 20 K. The same behavior was observed in   [31] . Moreover, it is possible to see that, as time advances, the ra  diation temperature does not change uniformly along the samples   length, but the distance between two curves corresponding to con  secutive instants is variable, and in some case the lines even cross   each other. This phenomenon is evidenced also in the pictures on   the right ( Figs. 6 and 7 (b)). The curves display an oscillatory trend   and tend to get closer and farther as time advances, testifying an   unsteady evolution of the radiated power which is different at each   axial location.   For a sound analysis of the material thermal behavior dur  ing arc-jet testing, infrared measurements on sample ZS10Y-LF   Fig. 8. Thermal images of sample ZS10Y-LF (a) before and (b) after the temperature   were taken as reference. During the tests, the ISQ5 pyrometer was   jump.   pointed towards the sample front face, whereas the IGAR6 pyrom  eter looked at the side surface. Moreover, the IR thermo-camera   positioning allowed to analyze both the front and the side surfaces   test, whereas, at temperatures over 2100 K, they started diverging,   facilitating the visualization and characterization of the tempera  and the front face reached a steady-state temperature around 500   ture jump phenomenon. Fig. 8 (a) and (b) show the thermal dis  K higher than the side.   tribution on the sample respectively just before and just after the   Thermo-camera and pyrometers measurements were compared   jump, during the maximum enthalpy step (H 0    18 MJ/kg). It is   to provide an estimation of the spectral emissivity in the near   =  evident that only the front part of the sample experienced a dra  infrared wavelength band, which is common to all instruments.   matic increase in temperature, whereas the rear body appeared to   Fig. 10 shows the temperature curves collected by pyrometers and   be almost unaffected by the thermal rise. The same trend is visible   in Fig. 9 , where the thermal histories recorded by the two pyrom  eters for the two surfaces of the sample are shown. It is evident   that the measurements matched well in the earliest phases of the   thermo-camera, on front (a) and on side (b) surfaces of sample   ZS10Y-LF, assuming a spectral emissivity    0.7. The non-perfect   ελ =  overlapping suggests a change in emissivity during heating. This  ελ , evaluated by   trend is quantified in Fig. 11 , where the value of   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   7   Fig. 9. Comparison of temperature measured by ISQ5 pyrometer (front surface) and   Fig. 11. Spectral emissivity in the near-infrared wavelength band on front and side   IGAR6 pyrometer (side surface) for sample ZS10Y-LF.   surface of sample ZS10Y-LF, versus test time.   matching the pyrometer and thermo-camera measurements as de  scribed in section 2.2 , is plotted versus test time for both the front   (black line) and side (gray line) surfaces of sample ZS10Y-LF. Spec  tral emissivity appears to increase in the earliest phases of the test,   from a value below 0.4 after 30 s (end of first enthalpy step) up to   over 0.7, after roughly 60 s (end of second enthalpy step). Then,  ελ gradually decreases during heating on both surfaces, attaining   a minimum value between 0.5 and 0.6 and then rapidly increas  ing again after the temperature jump up to almost 0.8. Finally, it   appears that, during the cooling phase, the emissivity rapidly de  creases on the front surface, while it is almost constant on the side.   Finally, Fig. 12 shows the trend the spectral emissivity of the sam  ple front surface versus temperature.   3.2. Post-test microstructure analysis   Fig. 12. Spectral emissivity in the near-infrared wavelength band for sample ZS10Y  LF, versus temperature.   In order to implement the numerical calculation and gain an   insight into the samples aerothermal behavior during arc-jet test  might be due to the thermal stress rising from coupling of ZrO 2   ing, post-test microstructures were analyzed by SEM and energy   with different pore size and volume fraction.   dispersive x-ray spectroscopy (EDS, INCA Energy 300, Oxford in  EDS elemental analysis of the cross section in the central area   struments, UK).   Comparative views of the external oxidized surface of the sam  of the model, Fig. 14 , evidenced an oxygen depth penetration of  μm in ZS5Y-SF and ZS10Y-SF, respectively, in   400 and about 550   ples with short fibers is reported in Fig. 13 . It can be appreci  agreement with the severity of the test conditions. Interesting to   ated that in both cases, the surfaces were fiber-free, ZS5Y-SF dis  note is Si distribution across the scale: in the less harsh condition,   played alternated areas with dense ZrO 2   or cracked and porous   ZS5Y-SF, Si accumulated towards the top surface and filled ZrO 2   ZrO 2  , Fig. 13 (a), whereas ZS10Y-SF was featured by diffused crack  scale in the sub-layer. In the hardest condition, ZS10Y-SF, no SiO 2    ing in a ZrO 2   layer with large pores which coated another one   based topping accumulated owing to its instant evaporation once   where finer pores could be seen, Fig. 13 (b). This structure suggests   the outermost surface once was achieved.   that partial oxide removal took place leaving uncoated the oxide   Magnified views of the cross section of the SF samples are   layer underneath. To note that XRD analysis on these surfaces con  shown in Fig. 15 . In both cases the unaffected core and oxide   firmed that the ZrO 2  scale maintained the tetragonal structure ow  scale had a strong and coherent interface, carbon fibers did not   ing to the stabilization by Y 2   O 3   . Therefore, cracking in ZS10Y-SF   survive in the ZrO 2   layer and silica-based accumulation was clear   Fig. 10. Comparison between temperature measured by pyrometers and thermo-camera (   ελ =   0.7) (a) on the front surface and (b) on the side surface, of sample ZS10Y-LF.   \\x0c', '8   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   Fig. 13. SEM images of the surface of the short fiber reinforced samples after arc-jet tests with enlarged views of microstructural details inset: a) ZS5Y-SF, b) ZS10Y-SF.   Fig. 14. SEM image and corresponding EDS elemental mapping of a) ZS5Y-SF, b) ZS10Y-SF samples taken from the central zone of the buttons.   only in ZS5Y-SF, Fig. 15 (a). In the case of ZS10Y-SF, despite the   posure to the supersonic plasma flow with EDS mapping of the   embrittlement of the sub-layer owing to the fiber consumption,   no spalling occurred, possibly owing to a stable oxide composi  main components and a zoom on the outermost oxide layer. The  μm, which ap  overall oxide layer had a thickness of around 360   tion which comprised ZrO 2   and a mixed Y-B-C-O phase. As for   proximately corresponds to the thickness of one fabric scale. The   the matrix composition containing 5 or 10 vol% SiC in ZS5Y-SF   oxide appeared more damaged and brittle in the center as com  and ZS10Y-SF, respectively, one could think that higher amount of   pared to the periphery where it was rather compact smooth. Below   SiC would provide better oxidation resistance, owing to the higher   this layer, the pristine microstructure appeared unaffected, with   source of protective silica glass [38] . However, under these exper  evidence of the 0 °/90 ° architecture of the fibers. EDS revealed com  imental conditions, SiC volume fraction seems not to play such a   plete boron removal in the region of oxygen penetration, whilst sil  relevant role, because most of the glass migrated to the surface   icon was slightly depleted in the subscale and accumulated close   (which could retard the advancement of the oxidation front) was   shear-transported away and vaporized at the ultra-high tempera  tures of 270 0-280 0 K.   to the surface. At higher magnification, the outermost layer com prised a roughly 50  μm-thick pure ZrO 2  region, which displayed an   irregular and highly porous structure. Below this, a silica-rich layer   Moving to the long-fiber UHTCMC, the outer surface presented   was present, which covered a SiC-depleted region, where voids   discontinuous zones of silica-based rich glass and a sort of ZrO 2   were left by Si outwards diffusion. Also in this case, there was no   sheath of the carbon fiber which were ablated away, Fig. 16 . Fig. 17   evidence of carbon fibers survival to the exposure to the arc-jet   shows SEM pictures of the ZS10Y-LF sample cross-section after ex  flow, all along the oxide thickness.   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   9   Fig. 15. SEM images of the cross sections of a) ZS5Y-SF, b) ZS10Y-SF samples showing the oxide architecture.   Fig. 16. SEM images of the surface of the long fiber reinforced sample after arc-jet tests, ZS10Y-LF, with enlarged views of microstructural details inset.   It can be stated that the long fiber configuration was very effec  tions to the phase changes occurred on the samples surfaces due   tive in limiting oxygen penetration across the material depth, pos  to ultra-high-temperature oxidation.   sibly thanks to the consumption of one carbon fabric sheet that   Since the exposed surface had no protective coating, surficial   locally enriched the environment in CO, thus contrasting oxygen   carbon fibers vaporized at relatively low temperature [39] , leav  advancement.   4. Discussion   ing a rough ZrB 2  -SiC surface. In agreement with ZrB 2  -SiC oxida  tion mechanism [ 40 , 41 ], at temperatures below 1400 K (where the   spectral emissivity is minimum, roughly 0.3), the sample surface   was covered by a boron-oxide glassy phase, as relevant SiC oxi  dation had not been triggered yet. Between 1500 and 1600 K, SiC   Several phenomena were observed during UHTCMC samples ex  underwent passive oxidation and the formation of a stable borosil  posure to arc-jet supersonic plasma flow. A comprehensive inter  icate glass (BSG) provided the highest oxidation protection to the   pretation could be proposed, relating the experimental observa  UHTC matrix, Fig. 18 (a). The maximum emissivity of around 0.75   \\x0c', '10   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   Fig. 17. SEM image of the cross section of the ZS10Y-LF sample after the arc-jet test with corresponding EDS elemental maps showing boron, oxygen and silicon distribution   across the profile. Bottom pictures show the outermost oxide layer at higher magnification.   Fig. 18. Sketch of LF UHTCMCs oxidation highlighting three different stages, a) formation of borosilicate glass and compact ZrO 2 layer, b) glass bubbling, development of   columnar ZrO 2 outer scale and progressive silica migration outwards from the subscale, c) exposed porous ZrO 2 layer after complete silica shear removal.   was reached at temperatures between 1600 and 1700 K, when con  teverde et al. [31] , the waves phenomenon was associated to the   sistent B 2   O 3   vaporization occurred and the glass was mostly com  shear-induced transportation of the glass along the specimen side   posed of SiO 2  . In the 210 0-230 0 K temperature range, which are   surface. The BSG layer is known to perform a protective action for   the maximum temperatures reached before the occurrence of the   the ceramic, acting as a barrier preventing further oxygen diffusion   temperature jump for all the samples, emissivity was then mini  and consistent material oxidation at intermediate-high tempera  mum, around 0.5. This is also the temperature range in which the   tures [43] . A combination of shear transportation and volatilization   waves-of-radiance phenomenon occurred, right before the temper  of the glassy phase left the underneath skeleton unprotected, ex  ature jump. Since at those temperatures the BSG layer was sup  posing the ZrO 2   grains directly to the supersonic flow [41] . Mean  posed to be mainly composed of silica, that has a melting point   while, the residual liquid phase being generated in the sub-scales   around 20 0 0 K [42] , we speculate that the oxide was completely   was not capable to prevent massive volatilization of the gaseous   liquid and therefore, in agreement with the discussion by Mon  products of SiC and carbon fibers oxidation (SiO, CO 2  , CO), whose   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   11   Fig. 19. Distributions of (a) Mach Number and (b) static temperature inside mixing chamber and nozzle, conditions corresponding to Step 7 (H 0   =   18 MJ/kg).   Fig. 20. Distributions of (a) static pressure, (b) static temperature, (c) O mass fraction and (d) N mass fraction around the sample, conditions corresponding to Step 7   =  (H 0    18 MJ/kg).   vapor pressure led to unsteady protrusion and bursting of liquid   in the assumption that, especially at the very high temperatures   bubbles, as observed in a previous test campaign [26] , Fig. 18 (b,c).   reached after the jump, the estimated value of spectral emissiv  This is coherent with the microstructural analysis presented in   ity is representative of the total emissivity along the whole wave  section 3.2 , showing that the outer oxide layer on all the samples   consisted of porous zirconia, after massive removal of both carbon   fibers and silica glass.   length spectrum (as most of the power is irradiated in the wave μm, where the value of  ελ was calculated   length band around 1   [48] ), it could be argued that the observed increase in the emissiv  At this point, some considerations can be done about the tem  ity after the jump should even favor an improved radiative heat   perature jump, observed for all the samples at a flow total en  dispersion. On the other hand, the general estimated emissivity   thalpy over 18 MJ/kg. This phenomenon has been observed by   trend is in agreement with the total emissivity measurements pre  several authors for SiC-containing UHTCs and C-SiC, but a widely   sented by Scatteia et al. [49] , while the increase in spectral emis  agreed interpretation is still lacking. The mechanisms proposed as   sivity with temperature, observed after the occurrence of the jump,   possible triggers for the jump include transition from passive to   is coherent with available data related to total emittance of ZrO 2  at   active oxidation of silicon carbide [ 44 , 45 ], triggering of catalytic re  ultra-high-temperature [50] , so a different reason is most likely to   combination of nitrogen atoms due to the presence of gaseous sil  be searched.   icon [46] , formation of cracks promoting oxygen diffusion to inner   The experimental evidence, with the temperature measurement   SiC particles and carbon fibers, resulting in carbon exothermic oxi  showing the jump being confined in a thin region in proximity of   dation and nitridation [ 44 , 47 ], surface modifications altering prop  the sample front surface, and the SEM images revealing a porous   erties of the samples such as emissivity and catalycity. All these   oxide layer mainly composed of zirconia, suggests instead that a   factors could lead to completely different surface heat flows even   twofold mechanism could be taken into account to justify the tem  under the same test conditions (same arc power, in the present   perature jump:   case). It is important to underline that, whatever the reason be  hind the considerable temperature increase, this was not just a   transitory phase, but led to a new condition which persisted for all   the remaining test duration. Indeed, the plots of Fig. 4 demonstrate   that a steady-state radiative equilibrium temperature was achieved,   not only during the maximum enthalpy step, but even during the   cooling procedure, in which the torch arc power is stepwise de  creased. In the cooling phase, the surface temperature measured   by the ISQ5 pyrometer was always higher than during the heating   a substantial reduction of the thermal conductivity in the oxide   layer, on which porosity itself can have a significant influence   [ 51 , 52 ];   an increased catalytic recombination efficiency due to a transi  atively low   tion from an oxide layer mainly covered by glassy silica (rel γ w [ 53 , 54 ]) to a scale primarily based on zirconia  (relatively high  γ w [55] ).   In the next section, this interpretation will be quantitatively   sequence, at all the enthalpy levels. Therefore, after the unsteady   supported by the outcomes of numerical simulations, which will   evolution corresponding to the trigger of the temperature jump, a   be compared with the experimental data.   new stable equilibrium condition of the heat balance through the   exposed surface of the sample must have been established, related   4.1. Numerical simulations   to variations of the heat flux contributions at the solid/fluid inter  face.   Computational Fluid Dynamic simulations were employed to   One possible trigger could be a reduced capability of the ma  get more detailed information about the evolution of the aero  terial to dissipate the incoming heat flux by radiation. However,   thermo-chemical flow field that develops in the facility and around   \\x0c', '12   S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   and to 0.8 after the jump. In order to match the temperature ax  ial profile before the jump, and based on experimental measure  ments performed within the frame of the project, a temperature   dependent thermal conductivity was considered, varying between  49 W/ m  \\x16K at room temperature to 47 W/m  \\x16K over 2273 K. Even   before the occurrence of the temperature jump, a certain amount   Fig. 21. Non-catalytic and fully catalytic cold wall heat flux profiles on the front   surface of the samples, conditions corresponding to Step 7 (H 0    18 MJ/kg).   ation of the surface chemistry, resulting in a complete removal   of catalytic recombination needed to be taken into account, with  \\x1610  −3 . This value is representative of   a catalytic efficiency    4   =  γ w   SiO 2   [53] , which is supposed to be the component with the high  est concentration in the BSG phase before the jump. As discussed   in Section 3.1 , at the earliest stages of the temperature jump phe  nomenon, a liquid phase is transported downstream by the su  personic flow, generating the waves-of-radiance phenomenon. It is   opinion of the authors that this is the trigger for an unsteady vari  of SiO 2   and carbon fibers from the sub-scale layer of the oxide   phase, which proceeds until a new radiative equilibrium condition   is reached, corresponding to the exposure of a highly porous ZrO 2   scale (see Fig. 17 ). As anticipated above, to justify the rise in tem  perature associated to the jump, and localized in the front part of   the sample, a double effect was considered: an increase in catalytic  activity (  γ w ), and a dramatic decrease in thermal conductivity in   the oxidized region (k ox ). An excellent agreement between numer \\x1610   ical and experimental results was obtained assuming   −2   =   7   γ w   (more than one order of magnitude higher than before the jump)   1 W/m  \\x16K, and considering an oxide thickness of 400  μm (based on the estimations made by the microstructural analy  and k ox   =  ses presented in Section 3.2 ). The comparison between numerical   and experimental temperature axial profiles is shown in Fig. 22 .   The results of these CFD simulations support the proposed in  terpretation for the temperature jump, demonstrating that, with   reasonable assumptions about the involved physical parameters,   it is possible to accurately reproduce the experimental behavior   of the UHTCMCs, providing a sufficiently solid explanation of the   temperature jump phenomenon.   =  =  the samples, and to provide possible interpretations for the ther  mal behavior of the materials.   Specifically, to investigate the phenomenon of the tempera  ture jump, the last step of the test on ZS10Y-LF was selected   as a reference case. First, the thermo-fluid-dynamic and chemical   field was simulated, employing the numerical models described in   Section 2.3 . Fig. 19 shows the Mach number and temperature dis  tributions inside the mixing chamber and supersonic nozzle of the   SPES wind tunnel, in the selected conditions (H 0    18 MJ/kg). The   hot jet coming from the axial torch is clearly distinguishable.   The conditions obtained at the nozzle exit section were used as   inputs for the simulation of the flow field around the test article.   Fig. 20 (a-d) show the distributions of pressure, temperature and   mass fractions of dissociated oxygen and nitrogen. Static pressure   and temperature contours evidence the structure of the flow field,   with a normal shockwave forming at a distance of 4-5 mm from   the sample, and the aforementioned quantities rising downstream   the shock to values about 90 0 0 Pa and 60 0 0 K respectively. It is   also clear that the level of dissociation of molecular species is con  siderably high, condition that, as discussed below, results in a rel  5. Conclusions   evant effect of surface catalycity on the heat fluxes. Fig. 21 shows   in fact the profiles of the cold wall heat flux on the front surface   An experimental campaign was carried out to characterize a   of the sample, for both non-catalytic (NC) and fully catalytic (FC)   new class of Ultra-High-Temperature Ceramic Matrix Composites   conditions. It is evident that the heat flux is more than double in   based on ZrB 2  -SiC-Y 2   O 3   matrix and different carbon fibers archi  the FC condition with respect to the NC, reaching values around 5   tectures, in an environment representative of atmospheric re-entry.   MW/m   2 .   Small sized samples were exposed to a supersonic flow of simu  The aero-thermo-chemical field was finally coupled to the ther  lated air in an arc-jet wind tunnel, at specific total enthalpies up   mal analysis of the sample, performing steady-state simulations   to 20 MJ/kg, at heat fluxes over 5 MW/m  2 (cold wall, fully cat  to match the temperature distribution evaluated by the thermo  alytic) and in a highly reactive chemical environment. All samples   camera before and after the temperature jump. The sample density   was set to 4300 kg/m  3 . The surface emissivity, based on the esti  demonstrated an excellent ablation resistance, with erosion rates  −4 mm/s. The surface temperature was mon on the order of 10   mation presented in Section 3.1 , was set to 0.5 before the jump,   itored by non-intrusive infrared equipment, including two-color   Fig. 22. Comparison between numerical and experimental temperature axial profiles, before (left) and after (right) temperature jump, for sample ZS10Y-LF.   \\x0c', 'S. Mungiguerra, G.D. Di Martino and A . Cecere et al. / International Journal of Heat and Mass Transfer 156 (2020) 119910   13   pyrometers and a thermo-camera. In all the tests, a spontaneous   temperature jump of several hundred degrees was observed at   constant flow conditions, with maximum equilibrium surface tem  peratures of 2690-2850 K. Only a thin layer in the front part of the   samples, directly exposed to the flow, experienced the jump, while   the rear material kept a much lower temperature (below 2200 K   even after the jump). A phenomenon of liquid phase transporta  tion along the side of the sample ( waves-of-radiance ) was observed   on the samples right before the onset of the jump. The combined   effort of experiment al activities and numerical simulations allowed   proposing a novel and comprehensive interpretation for the jump,   based on a twofold mechanism affecting the heat transfer to the   material, associated to the formation of a porous ZrO 2  layer on the   external surface of the sample after complete removal of the liquid   SiO 2  glassy phase: an increase in the catalytic activity and a strong   reduction in thermal conductivity in the oxidized region. Despite   the jump, the analyzed ceramics displayed promising performance,   with an excellent ablation resistance at 270 0-280 0 K and the capa  bility of the oxide phase to shelter the underneath material, keep  ing it at acceptable temperatures.   CRediT authorship contribution statement   Stefano Mungiguerra: Conceptualization, Methodology, Valida  tion, Formal analysis, Investigation, Data curation, Writing origi  nal draft, Visualization. Giuseppe D. Di Martino: Methodology, In  vestigation, Writing review & editing. Anselmo Cecere: Investi  gation, Writing review & editing. Raffaele Savino: Conceptual  ization, Methodology, Supervision, Funding acquisition. Luca Zoli:   Methodology, Investigation, Resources, Writing review & editing.   Laura Silvestroni: Methodology, Investigation, Resources, Writing   review & editing. Diletta Sciti: Supervision, Funding acquisition,   Project administration, Writing review & editing.   Acknowledgments   This study has received funding by the European Union’s Hori  zon2020 research and innovation programme under the research   project C  3 HARME with Grant Agreement No. 685594 .   Supplementary materials   Supplementary material associated with this article can be   found, in the online version, at doi:10.1016/j.ijheatmasstransfer.   2020.119910 .   References   [1] W.G. Fahrenholtz , G.E. 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Klemens , Thermal conductivity of dense and   Boca Raton, FL, 2011 .   porous Yttria-stabilized zirconia, J. Mater. Sci. 36 (2001) 3003-3010 .   [43] S. Gangireddy , S. Nanna Karlsdottir , J.W. Halloran , Liquid oxide flow   [52] B. Nait-Ali , K. Haberko , H. Vesteghem , J. Absi , D.S. Smith , Thermal conductivity   during   oxidation   of   zirconium   diboride-silicon   carbide   ultra   high   of highly porous zirconia, J. Eur. Ceram. Soc. 26 (2006) 3567-3574 .   temperature ceramics, Key Eng. Mater. 434 (2010) 144-148 435doi:   [53] L. Bedra , M. Balat-Pichelin , Comparative modeling study and experimental re  10.4028/www.scientific.net/KEM.434-435.144 .   sults of atomic oxygen recombination on silica-based surfaces at high-temper  [44] I. Sakraker , C.O. Asma , Experimental investigation of passive/active oxidation   ature, Aerosp. Sci. Technol. 9 (2005) 318-328 .   behavior of SiC based ceramic thermal protection materials exposed to high   [54] Y. Yang , I. Kim , G. 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},{
  "_id": 280,
  "PDF": "Ultra-refractory ceramics the use of sintering aids to obtain microstructure control and properties improvement.pdf",
  "Text": "['Key Engineering Materials Vols 264-268 (2004) pp 787-792 © (2004) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.264-268.787  Online: 2004-05-15  All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, www.ttp.net. (ID: 131.173.17.71, UB Osnabrück, Osnabrueck, Germany-26/05/15,15:47:56)  Ultra-Refractory Ceramics: the Use of Sintering Aids to Obtain Microstructure Control and Properties Improvement  A. Bellosi and F. Monteverde  CNR-ISTEC, Institute of Science and Technology for Ceramics,  Via Granarolo 64, 48018, Faenza, Italy  Keywords:   UHTC, zirconium diboride, hafnium diboride, silicon carbide, ceramic composites,                       sintering,  microstructure, mechanical properties, oxidation resistance  Abstract. The present study focuses on innovative processing for the densification of refractory diborides (ZrB2, HfB2) based ceramics, suitable to fabricate UHTC components for structural applications and thermal protection systems. The addition of small quantities (2-5 vol%) of sintering aids allows to overcome the intrinsic low sinterability of zirconium diboride and to improve properties through the microstructural refinement. The introduction of second refractory phases, for instance 20 vol% of SiC in the diboride matrices improves strength, toughness and oxidation resistance. Outstanding properties were measured: fracture toughness 3-5 MPam1/2, hardness 9-20 GPa, elastic modulus 350-420 GPa, flexural strength 300 to 750 MPa at R.T and up to 300 MPa at 1500°C. The resistance to oxidation of different materials was compared.   Introduction  Refractory diborides are materials of particular interest because of high melting point, high electrical and thermal conductivity, chemical inertness against molten metals or non basic slag, superb thermal shock resistance [1]. These properties make it an attractive candidate for high temperature applications where corrosion-wear-oxidation resistance are demanded.  Applications of this ceramic as a refractory in foundry or electrical devices (heaters, igniters) are currently in use. Other applications have been recently proposed for thermal protection systems in aerospace [2-4]. These ceramics can be electrical discharge machined to produce complex shaped parts, due to their low electrical resistivity ( 10 µΩcm). The relatively low strength and fracture toughness are still obstacles to a wider diffusion of ceramic borides. In order to improve properties, control of the processing parameters and of the final microstructure are required. Pressure assisted sintering techniques are necessary to densify pure diborides-based ceramics [1, 5], generally at temperatures higher than 1950°C that cause microstructure coarsening. The addition of sintering aids represents a change to overcome the intrinsic low sinterability [1, 6-8] and to refine the microstructure. However, in some cases, the grain boundary phases deriving from the sintering aids deteriorate the properties [6, 8-12], therefore the control of the properties of borides can be achieved introducing second phases [12, 14, 15].  The present results highlight the improvement in sinterability, microstructure and properties of ZrB2 and HfB2 ceramics produced with the addition of different sintering aids, either metal (Ni) or ceramic (Si3N4 and AlN), compared to an additive-free material. Besides, the influence on the properties of the addition of 20vol% of SiC as reinforcing phase is evaluated.   Compositions, processing and microstructure of ultra-refractory ceramics   Several ZrB2and HfB2-based ceramics were produced starting from commercial powders of ZrB2, HfB2, SiC, AlN, Si3N4 and Ni [12, 13, 18]. Hot pressing conditions and final densities are shown in Table 1. The values of the final relative density of the hot pressed materials evidence the relevant role of the sintering aids. In fact pure ZrB2 (material 1) has a density remarkably lower than that of the other doped compositions. As described in details in previous works [12, 13, 16] the\\x0c', '788  Euro Ceramics VIII  densification behaviour supports the hypothesis that a liquid phase forms during hot pressing, at temperatures that depend on the reactions occurring among the additives and the species present in the starting ZrB2 or HfB2 powders, including B2O3 and impurities. Such liquid phase has two main effects: it favours powder particle rearrangement and drives the activity and diffusion of the atomic species involved in the densification. In fact,  during sintering, dissolution-diffusion-repricipitation mechanisms allow the powder compact to densify faster at lower temperature than in the absence of liquid phase. Depending on the type and amount of the additives, the reactivity among the involved phases and the phenomena governing the  Table 1. Compositions, hot pressing temperature (T) and time (t), ( P=30MPa),  final relative density (r.d.) of the materials  ------------------------------------------------------Material (vol%)    T (°C)   t(min) r.d.% ------------------------------------------------------1  ZrB2            1870        30    87 2  ZrB2+4 Ni         1850        30    98 3  ZrB2+5 Si3N4            1700        10    98 4  ZrB2+5 AlN         1850        30    92 5  ZrB2+19 SiC+s.a.*    1850          10    98 6  HfB2+19 SiC+s.a.*    1850          20   98 ------------------------------------------------------* s.a.=sintering aids matter transfer during densification are controlled by different factors, that determine either the sintering rate or the microstructure evolution.  In the case of the additive free material 1, grain growth is associated with retarded densification due to surface oxide (B2O3) that hinder diffusion mechanisms (Fig. 1). It is well known that oxygen contamination limits sinterability of borides due to evaporation/condensation kinetics.  On the contrary, ZrB2 containing Ni (material 2) reached a final density of 6.05 g/cm3 which is about the 98% of the theoretical density. ZrB2 grains are quite homogeneous in shape and size. Ni-rich grain boundary phases are located mainly at triple points (Fig. 2). Evaporation phenomena of ZrB2 was also strongly limited due to the presence of liquid phase (that forms at about 1260°C [12]), thus exaggerated grain growth was retarded. Being dependent upon the wetting behaviour, which is influenced by the surface oxidation of ZrB2 particles, some residual pores, partially filled of the liquid phase, remained at the triple junctions. The existence of traces of crystalline Ni2B and of low amounts of reaction products in the Zr-B-O, Zr-O or Zr-Ni-O systems as grain boundary phases confirms some degree of reactivity among the different elements during hot pressing. Ni reacts with the oxygen present as impurity in the powders; then exchange reactions with ZrB2 give Ni2B and ZrO2. As a confirmation, small ZrO2 particles formed, mostly within ZrB2 grains. Zr-O compounds were also present in round inclusions distributed mainly at grain boundaries (Fig. 2). Material 3 (Si3N4-doped ZrB2) has a relative density of about 98%. The microstructure is rather regular and fine (mean grain size about 3 µm), the residual porosity is very scarce.  Figure 1. Polished section of undoped ZrB2 Figure 2. Polished section of ZrB2-Ni      5 µ5 µ5 µ5 µm  10 µµµµm \\x0c', 'Key Engineering Materials Vols. 264-268  789  Grain boundaries appear generally flat, with secondary phases mainly concentrated at the triple points (Fig.3). Limited amounts of BN, ZrO2 and a glassy phase in the system B-N-O-Zr-Si were found. These phases formed during hot pressing by reaction among silicon nitride, boron and silicon oxide present on ZrB2 and Si3N4 particles, respectively, and ZrB2 itself. In fact, ZrB2 and Si3N4 cannot coexist under the adopted processing condition. The liquid phase formed at about 1580°C. The involved reactions [13] cause a complete disappearance of Si3N4 and produce various compounds that, after cooling, remain entrapped, mainly at triple points among the grains.  The addition of AlN as sintering aid in material 4 caused the partial removal of the B2O3 oxide layer that covers the ZrB2 particles, with formation of reaction products like Al2O3 and BN. In spite of the higher processing temperature, the lower relative density (92%), compared to material 3, suggests a minor reactivity among the involved compounds when Si3N4 is substituted by AlN. However, in spite of the rather low density, the remarkable feature of ZrB2 with the addition of AlN is the absence of glassy phases (Fig. 4), being the grain boundary phases constituted by refractory compounds like Al2O3, BN and residual AlN.   Fig. 3. Polished surface of ZrB2-Si3N4 Fig. 4. Polished surface of ZrB2-AlN      The ZrB2-SiC and HfB2-SiC composites (material 5 and 6, respectively) have very fine and homogeneous microstructures (Fig. 5, 6) [18]. The SiC particles present a clustered distribution within in the diboride matrix and are often in contact with the products of the sintering process: ZrO2 (or HfO2), BN, glassy phase containing Si, Zr (or Hf) and other impurity cations. Such microstructural features confirm, also in this case, that the composites densify by the aid of a liquid phase, having composition, melting point and viscosity dependent on the sintering aids [16].   Figure 5.Polished surface of ZrB2-SiC composite  Figure 6. Polished surface of HfB2-SiC composite  1 µµµµm  5 µµµµm  4 µµµµm \\x0c', '790  Euro Ceramics VIII  Mechanical Properties  Table 2. Physical and mechanical properties of the various samples   Material α (°C-1) Hv1.0 (GPa) E (GPa) KIc (MPa√m) σ (MPa) R.T. σ (MPa) 1000°C σ (MPa) 1200°C σ (MPa) 1400°C σ (MPa) 1500°C 1 7.5 8.7±.4 346 2.4±0.2 350±40 320±60 313±10 219±5 2 7.5 14.4±.8 496 3.4±0.4 370±25 320±20 15±5 3 13.4±.6 419 3.7±0.1 595±90 400±20 240±30 4 7.65 9.5±.5 407 3.1±0.5 580±80 290±20 200±5 5 7.5 14.2±.6 421 4.5±0.1 730±100 430±110 250±15   6 7.3 20.4±.6 4.3±0.1 560±100 380±50 280±20 150±5 α: Linear thermal expansion coefficient (20-1300°C); Hv1.0: microhardness (measured by a Vickers indenter, load 9.81 N) E: Young’s modulus (resonance frequency method); KIc fracture toughness (chevron notch method on 25.0x2.5x2.0mm3 test bars); σ: 4-pt flexural strength  (from R.T.up to 1500°C in air on 25.0x2.5x2.0 mm3 test bars, inner span 20 mm, outer span 40 mm, crosshead speed 0.5 mm/min)  Young’s modulus (E), microhardness (Hv1.0): No reference values were found in literature for comparison. Residual porosity, grain size and secondary phases influence these properties. Fracture toughness (KIc). The addition of SiC as a second phase in sample 5 and 6 introduced a change in the fracture mode, from intergranular to mixed inter/transgranular mode. The increase in fracture toughness of the two composites derives from a strengthening of the grain boundaries due to the SiC particles. Previous results [16] revealed that the ZrB2-SiC interfaces are stronger than ZrB2-ZrB2 interfaces, so we suppose that the same mechanisms are active also for HfB2-SiC. Flexural strength (σ): the R.T. values are higher than those found in literature due to the refined microstructure and the control of the defect population. Material 1 is penalized by the residual porosity. Material 2 has good values and low size defects. Materials 3, 4, 5 and 6 reveal outstanding strength due to the composition design and the consequent microstructure. High temperature strength is related to the characteristics of the grain boundary phases. When a metallic sintering aid was used (material 2) the softening of the grain boundary phases at T>800°C becomes detrimental. Material 3 and 5 evidence a slight strength decrease at T >1000°C. The strength at 1200°C (250 MPa), although it is about 60% of that at R.T., is high for this class of materials. Tests carried out up to 1500°C on materials 4 and 6 show excellent results. The strength decrease is mainly due to softening of glassy intergranular phases. Additional effects on strength at high temperature are related to the surface modification for oxidation.  Surface degradation for oxidation   The oxidation resistance of diboride ceramics depends on the microstructural characteristics. Specifically, the addition of SiC particles represents the key factor that makes the tested composites ZrB2-SiC and HfB2-SiC (samples 5 and 6) consistently more resistant to oxidation than the monolithic diborides. The weight gain for oxidation, during heating up to 1350°C, Fig. 7, is relative to two monolithic ZrB2 materials and one composite.Taking into consideration the phases present in  the as-sintered materials, both  condensed  and  gaseous  oxidation  products  can  form  during 01234567890200400600800100012001400T (°C)w (mg/cm2)3) ZrB2-Si3N4 1) ZrB25) ZrB2-SiCFigure 7. Weight change vs. temperature \\x0c', 'Key Engineering Materials Vols. 264-268  791   oxidation. The main reactions describing the oxidation process, which involve , either mass gain  ZrB2 + 5/2O2 = ZrO2 + B2O3            (1) 2BN + 3/2 O2 = B2O3 + N2               (2) SiC+3/2 O2 = SiO2 + CO                 (3) or mass loss B2O3 (l) = B2O3 (g)                              (4)                                                 SiO2(s) + CO(g) = SiO(g) + CO2(g)   (5) are simultaneously active. Therefore the interpretation of the weight change data has to take into consideration the formation of volatile products. The main reaction products from the oxidation of  ZrB2, which constitutes the basic matrix of the tested materials, are zirconia and amorphous boric acid (B2O3). Actually the latter compound has an unusual low melting point (450°C) and an high vapour pressure: at high temperature it vaporizes. The rapid increase in weight gain of material 1 (Fig. 7) reflects the weak ability of B2O3 to hinder intensive oxidation. On the other hand, the oxidation layer formed by ZrO2 has a semi protective action. Nonetheless, grain boundary silicon-containing phases, although in very limited amounts like in material 3, promote the formation of a glass on the material surface, that improves the resistance to oxidation because it acts as barrier against the inward diffusion of oxygen, more efficiently than B2O3 and ZrO2 do. The addition of SiC (material 5 and 6) enhances the production of a silica rich glassy phase on the top of the sample that ensures an almost continuous coating of the sample surface, either with an outer glassy layer or with a glass/ZrO2 (or glass/HfO2) composite sublayer, both inhibiting the inward access of oxygen.  Therefore, the presence of silicon carbide as particulate second phase is certainly responsible for the noticeable improvement in the oxidation resistance. In the very first oxidation step the major part of SiC particles acts as an effective obstacle against the inward diffusion of oxygen and, at the same time, their oxidation improves the formation of a silica-enriched glass. This melt, fluid at the testing temperature and characterized by a restricted permeability to oxygen, seals quickly the external surfaces of the test piece and short-circuit paths for the incoming oxygen (i.e residual porosity, cracks) towards easily accessible ZrB2 (or HfB2) grains and/or grain boundary phases. The evidence of residual unoxidized SiC particles within the reaction scale settled indeed the fundamental role of such phase in suppressing the advance of the oxidation attack.     Oxidation tests carried out at 1600°C in air for 15 minutes confirmed the formation of a surface glassy coating as the key point for the improvement of the oxidation resistance also at very high temperatures. The cross sections of the oxidized samples after this test are shown in Figs. 8a-c and 9a-c. It is evident that in the materials 1 and 4 the transformation of ZrB2 into porous ZrO2 enters towards the inner bulk for a thickness of about 0.5 mm. On the contrary, in the samples 3, 5 and 6, a glassy surface layer, about 10 µm thick, protects the bulk from the inward diffusion of oxygen. A relatively thin (about 50 µm)  subsurface layer is interested by oxidation.    Undoped ZrB2 (sample 1)  ZrB2+AlN (sample 4)  Composite ZrB2-SiC (sample 5)  Figure 8. Cross sections of samples oxidized at 1600°C for 15 min, showing the depth of oxidation     0.5 mm \\x0c', '792  Euro Ceramics VIII     Figure 9. Details of the cross sections after oxidation at 1600°C; upperside: the protective surface glass.  In the case of material 3 (Si3N4-doped ZrB2) the small amount of intergranular glassy phase is enough to produce a surface silica-based coating. However, it is not continuous and well adherent to the bulk like the glassy layer formed in the composites, containing silicon carbide (materials 5 and 6). The scarce oxidation resistance of samples 1 and 4, that do not have grain boundary phases, derives from the combination of two factors: the partially protective nature of the surface oxide and the presence of residual porosity, that provide easy pathes for oxygen to enter towards the bulk and react with the inner zirconium diboride grains.   Conclusions  Highly dense refractory-based ceramics can be sintered at temperatures in the range 1700-1900°C, with the addition of proper sintering aids. Materials exhibiting excellent mechanical properties up to 1500°C have been obtained with the addition of aluminium nitride as sintering aid. Enhanced oxidation resistance was observed in ZrB2-SiC and HfB2-SiC composites due to the formation of a surface glassy layer protecting the bulk from the contact with oxygen. The composition design is the key factor for the thermal stability of these materials at very high temperatures.  Acknowledgement This work was supported by ASI (Italian Space Agency), Project NEIN700257. The Authors thank D. Dalle Fabbriche, C. Melandri and S. Guicciardi for their helpful collaboration.  References  1. C. Mroz: Am. Ceram. Soc. Bull.  Vol. 74, N°6, (1995), p. 165. 2. I. M. Low and R. McPherson: Mat. Sci. Lett.Vol. 8, (1989), p. 1288. 3. K. Upadhya, J. M. Yang and W.P. Hoffmann: Am. Ceram. Soc. Bull. Vol. 58, (1997), p.51. 4. G.P. Van de Goor and K. Berroth, in Advanced Multilayered and Fibre-Reinforced Composites, (Kluwer Academic Publ, 1998) p.311. 5. R.M. Hayami, H. Iwasa. and M. Kinoshita: Yogyo-Kyokai-Shi Vol. 88, (1978) p. 352. 6. A. Bellosi and F. Monteverde, in Key Engineering Materials Vols. 175-176, 130-138, 2000. 7. S-K Woo, I-S Han, H-S Kang. and I-H Yang: J. Korean Ceram. Soc. Vol. 33, 259, 1996. 8. A. Bellosi, F. Monteverde F.and C. Meandri: J. Mat. Proc&Manuf. Sci. Vol. 9, (2000)p. 156. 9. D.N. Øvrebø and F.L Riley, in Sixth ECerS Conference, Extended Abstract,  Vol 2, 2-11, 1999. 10. M.Rahman , C.C.Wang, S.A Akbar and C. Mroz: J. Am. Ceram. Soc. Vol. 78 (1995)p. 1380. 11. J.J. Schuldies and J.A. Branch: Ceram. Indust, Vol. 138, (1992) p. 43. 12. F. Monteverde, A. Bellosi and S. Guicciardi: J. Eur. Ceram. Soc. Vol. 22, 3, (2002) p. 278. 13. F. Monteverde and A. Bellosi: Scripta Materialia, Vol. 46, (2002) p. 223. 14. C.C. Sorrell, V. S. Stubican and R. C. Bradt : J. Am. Ceram. Soc. Vol. 69 (1986) p. 317. 15. T. Ogata, T. Mori, K.Nakamura, K. Kobaiashi and H. Kuwajima: Advanced Structural Inorganic Composites, (Elsevier Science Publishers, B.V, 1991).p.199. 16. F. Monteverde, S. Guicciardi and A. Bellosi: Mat. Sci Eng. A, Vol. 346 (2003) p. 310  25 µµµµm ZrB2+SiC (sample 5) HfB2+SiC (sample 6) ZrB2+Si3N4(sample 3) \\x0c', 'Euro Ceramics VIII   10.4028/www.scientific.net/KEM.264-268   Ultra-Refractory Ceramics: The Use of Sintering Aids to Obtain Microstructure Control and Properties Improvement   10.4028/www.scientific.net/KEM.264-268.787   DOI References  [13] F. Monteverde and A. Bellosi: Scripta Materialia, Vol. 46, (2002) p. 223.  doi:10.1016/S1359-6462(01)01229-5  [14] C.C. Sorrell, V. S. Stubican and R. C. Bradt : J. Am. Ceram. Soc. Vol. 69 (1986) p. 317.  doi:10.1111/j.1151-2916.1986.tb04739.x         \\x0c']"
},{
  "_id": 281,
  "PDF": "Ultrahigh-Temperature Ceramic Based on ZrB2–SiC- Preparation and Main Properties.pdf",
  "Text": "['ULTRAHIGH-TEMPERATURE CERAMIC BASED ON ZrB2-SiC:  PREPARATION AND MAIN PROPERTIES  P. S. Sokolov,1,2 A. V. Arakcheev,1 I. L. Mikhal’chik,1 L. A. Plyasunkova,1  A. V. Tkachev,1 S. A. Anuchin,1 M. N. Kordo,1 A. V. Lanin,1 A. O. Zabezhailov,1  I. Yu. Kelina,1,3 and M. Yu. Rusin1  Translated from Novye Ogneupory, No. 1, pp. 33 - 39, January, 2017.  Original article submitted September 22, 2016.  Hot pressing of commercially available powders is used to prepare dense ceramic based on ZrB2-SiC with dif ferent  additives  (Si3N4, TiSi2, ZrSi2, MoSi2). The main physicomechanical properties  are measured in  macro-specimens: ultimate strength in static three-point bending at room temperature 400 - 600 MPa, Vickers  microhardness up to 15 GPa, and critical stress intensity factor up to 5.9 MPa·m1/2. Average LTEC,  thermal  diffusivity, thermal conductivity, and oxidation resistance over a wide temperature range are determined. Al together the ceramic properties obtained are at the level of published indices.  Keywords: ultra-high temperature ceramic materials (UHTCM), ceramic based on ZrB2-SiC, hot pressing  (HP), zirconium diboride, silicon carbide.  INTRODUCTION  At  the start of  the 1960s  in the USSR and then in the  USA systematic and extensive research was conducted for  the properties of  refractories  (Tm > 1600°C) and ultrahigh temperature ceramic materials (UHTCM, Tm > 3000°C) such  as  borides,  nitrides,  carbides,  silicides,  phosphides,  etc.,  [1 - 3]. Apart from refractoriness many of these compounds  exhibit other unique properties: good chemical resistance to  various corrosive media, heat  resistance and refractoriness.  Research conducted for finding and selecting technology for  promising  materials  has  been  undertaken  by  leading  branches of industry; chemical, atomic, space, metallurgical,  microelectronics, and also for creating hypersonic spacecraft  (HS). Within the USSR this class of work was concentrated  in  the  Institute  of Materials  Science  Problems  (Kiev,  Ukraine), and in a number of enterprises of the atomic indus try. Production of UHTCM was set up in the Donetsk Chemi cal Reagent Plant. After the break-up of the USSR work on  this theme was wound up and almost ceased as a whole.  At  the  start of  the 2000s  throughout  the whole world  there was regeneration of interest in UHTCM and their tech nology [3 - 5]. Research emerged at a new quality level, new  analytical and synthetic equipment was used, extensive inter national  integration of scientists was developed, and consor tia of organizations have been created for different forms of  ownership,  and most  important  financial  support  for  this  work has reappeared. Overseas financing proceeds along the  lines of  the  largest  state  scientific  funds  and government  agencies for space or military themes. In Russia this work of  important science and technology is only carried out in small  collectives of academic institutes at an initiative level due to  internal  reserves or as a result of  financial support  through  RFFI lines [5 - 9].  Recently according to data of  the Google. Scholar sys tem explosive growth is observed in the amount of research  work [3 - 5] devoted to the theme of UHTCM based on ZrB2  and HfB2; from tens of works at the start of the 2000s to sev eral hundred currently. This is connected with two factors:  with activation of work in traditional centers concerned with  the search and study of materials for extreme applications (in  Europe, Japan, and USA), and with significant expansion of  the geography of research, and appearance of new centers of  interest  in other countries (China,  India,  Iran, Turkey, etc.).  In the last  two to three years a clear shift has been observed  Refractories and Industrial Ceramics  Vol. 58, No. 1, May, 2017  46  1083-4877/17/05801-0046 © 2017 Springer Science+Business Media New York  1  AO ONPP A. G. Romashin Tekhnologiya, Obnisk, Kaluga Re gion, Russia.  2  sokolov-petr@yandex.ru  3  kelina@technologiya.ru  DOI 10.1007/s11148-017-0052-9  \\x0c', 'in the theme of published articles from a purely academic to  a practical plane. For example, production questions are be ing discussed about effective molding of “starting materials”  (including  use  of  additive  technology), machining  of  ce ramic, welding  and  other  rational methods  for  joining  UHTCM.  Researchers  are  drawn  to  the  clear  advantages  of  UHTCM based on ZrB2 and HfB2 compared with traditional  refractory materials: working temperature above 1800°C and  a metal nature of  the bond, providing good electrical and  thermal conductivity, splendid thermal shock and heat resis tance, moderate thermal expansion, high hardness, and ac ceptable crack and wear  resistance. At  the same time,  there  are some retarding factors  still preventing universal exten sion of work with UHTCM. Primarily is  a  shortage  and  expensiveness of starting raw material.4 Another  important  factor is the complexity of the equipment section connected  with  preparing  powders  and  dense  ceramic  (massive  high-temperature vacuum furnaces, hot pressing units, etc.)  and also the difficulty of testing ceramic properties.  In view of the technical complexity of preparing mono lithic ceramic from pure ZrB2 and HfB2 powders it is normal  to used SiC as an additive reducing sintering temperature.  Silicon  carbide  (in  an  amount  of  20 - 30 vol.%) within  UHTCM composition fulfils several  important  functions:  it  is a grain growth inhibitor for the target phase, there is an in crease in oxidation resistance, strength, and thermal conduc tivity. Transition metal  carbides  and silicides  are used as  sintering additives [3, 4].  Previously  in  cooperation with GNTs FGUP Keldysh  Center spark plasma sintering (SPS) has been used to prepare  dense ceramic (relative density 92 - 93%) of compositions  ZrB2 - 20 wt.% SiC and ZrB2 - 5 wt.% Si3N4 [7] in the form  of  tablets up to 20 mm in diameter.  Indentation into these  tablets was used to measure microhardness  (HV = 10 - 18  GPa)  and  critical  stress  intensity  factor  (KIc = 3.5 - 5.9  MPa·m1/2) with a load of 2 kg. The small size of tablets and  the limited working volume of  the spark plasma sintering,  does not make it possible to measure other ceramic structural  properties. Therefore  in  the  next  step  subsequent  experi ments were  conducted in a hot-pressing (HP) unit with a  larger working volume.  Currently the research object  is ceramic of the composi tion ZrB2 - 30 vol.% SiC with different  additions  (Si3N4,  TiSi2, ZrSi2, MoSi2), and experimental data have been ob tained for the main thermophysical and strength properties of  this ceramic over a wide temperature range.  EXPERIMENTAL SECTION  Starting materials  Finely dispersed ZrB2 and titanium, zirconium, and mo lybdenum disilicide pure grade powders produced domesti cally were used. The second main ceramic component was  selected  as  commercial  grade  abrasive SiC powder. The  microstructure and composition of powders were studied by  x-ray phase  analysis  (XPA), optical,  and electron micros copy.  Rapid  analysis  of  the  chemical  composition  of  ZrB2 was performed by x-ray microanalysis (semi-quantita tive XMRA). The impurity content of silicon carbide powder  was determined by quantitative chemical analysis. Powder  specific surface Ssp was measured by gas filtration [11] in a  PSKh-9  instrument. Powder  pycnometric  density  r  p was  evaluated in an AccuPyk II 1340 helium pycnometer. Parti cle size distribution was obtained by laser diffraction in a  Fritsch Analysette  22 Microtec  plus  instrument. Measure ments  were made  in  distilled  water  with  addition  of  surfactant by ultrasonic treatment. Calculation of  the distri bution was made by Mi theory. The main properties of start ing raw materials are provided in Table 1.  Synthesis  Starting powders were screened through a brass screen  and working  fractions were  collected  (<40 mm for ZrB2,  <50 mm for  silicides). Silicon carbide was not previously  screened. Then wet powder mixing was performed in a labo ratory planetary mill with balls and milling vessel of silicon  nitride for 24 h with a drum rotation frequency of 150 rpm.  Then the charge was extracted from milling vessel, dried and  rubbed through a steel screen with a cell size of 200 mm. All  operations were performed in air. The mixture prepared in  this way was placed in an HP unit of original construction  with a  force up to 10 tons. Compaction and simultaneous  sintering was carried out  in graphite molds,  lubricated with  boron nitride. Compaction pressure was about 20 MPa, max imum temperature 1800 - 1900°C, and holding at  the maxi mum temperature for 30 min.  Specimens Nos. 1 and 2 were prepared from ZrB2 pow der produced by OOO KTM (ZrO2 < 3 wt.%), and the rest of  the specimens were prepared from ZrB2 produced by OOO  Ékos-Ural  (ZrO2 ~ 7 wt.%).  In specimen No. 1 SiCw (cubic  b-SiC) fibers were used [7]. Specimens Nos. 3, 4, 5a, and 6a  were prepared from SiC of  the Zaporozh’e Abrasive Com bine, and specimens Nos. 5b and 6b were prepared from SiC  from the Volga Abrasive Plant  (see Table 1). Specimen 6a  was not dense and synthesis was established with achieve ment of a temperature of about 1700°C with holding.  Ultrahigh-Temperature Ceramic Based on ZrB2-SiC: Preparation and Main Properties  47  4  Any compounds of boron, zirconium, and hafnium are classified  as strategic raw materials, on which there are export-import limi tations. According to [10] within Russia currently there is almost  complete absence of in-house production of zirconia refractories.  \\x0c', 'Ceramic sample preparation and analysis  Actual density r a, open porosity P, and water absorption  w were measured by hydrostatic weighing using distilled wa ter at 20°C. Theoretical density r  t was calculated according  to mixing,  taking for the density of ZrB2, SiC, Si3N4, TiSi2,  ZrSi2,  and MoSi2,  6.12,  3.21,  3.21,  4.02,  4.88,  and  6.26 g/cm3 respectively.  Impurity content  (mainly ZrO2)  in  calculating r  t was not considered. The relative density was  calculated from the ratio r a/r  t.  In order to measure ultimate strength in static three-point  bending (s ben) at different  temperatures specimens were cut  into bars with a size of 7.0 ´ 7.0 ´ 80.0 mm and then ground  to roughness Ra ~ 0.3 mm with a diamond tool.5 Chamfers  were removed from all bars in order to minimise the effect of  stress  concentration. Bend  testing was  carried  out  by  a  three-point bend method according to GOST 3309 with a  loading rate of 1.5 mm/min and distance between supports of  50 mm using UMM-5  9024 DP 100/1500  and P-0.5 ma chines. Testing at elevated temperature (1200 - 1400°C)  in  air was carried out by a scheme: a specimen was placed in a  heated furnace and remained within it for 10 - 15 min in or der  to equalize temperature and then it was  tested. Micro hardness was measured at 20°C by the Vickers method (load  from 0.05 to 3 kg,  indentation time 10 sec)  in polished sur faces  (Ra ~ 0,03 mm) using a Stuers DuraScan 50 hardness  meter. Measurements were conducted by three-point bending  at each load, and K1c was evaluated by the Palmquist method  through the length of radial cracks formed about an indenta tion by a diamond pyramid with a load of 3 kg. Ceramic  phase composition was determined in a polished surface by  means of a DRON-6 x-ray diffractometer (Cu Ká-radiation).  Microstructure was  studied by optical and electron micro scopes with an XMRA attachment.  Linear thermal expansion coefficient (LTEC) was deter mined in specimens with a size of 4.0 ´ 4.0 ´ 50.0 mm in air  in  the  range  20 - 1100/1300°C in  a Netzsch DIL 402C  dilatometer previously calibrated using a standard of alumi num oxide  ceramic  of  similar  size. LTEC measurements  were performed according to GOST 10978-2-14; data were  presented as  the  average LTEC (a  20-t)  in the  range  from  20°C to t. Thermal conductivity l was calculated as the prod uct of r  p, heat content cp, and thermal diffusivity a. The con tribution of thermal expansion was considered in measuring  density. A laser flash method was used to measure a in speci mens  10 ´ 10 ´ 2.5 mm in  size  at  up  to  700/100°C in  a  Netzsch LFA 457 instrument in air. Differential scanning cal orimetry (DSC) was used to measure cp  in a Netzsch STA  DSC 204F1 instrument  in a stream of nitrogen with linear  extrapolation to high temperature. Specimen oxidation resis tance was  tested in a Netzsch STA 449F1 instrument  in a  stream of commercial purity nitrogen at up to 1500°C and  also in a stationary atmosphere in a resistance muffle furnace  in  the  range  1300 - 1800°C. UHTCM specimens were  heated at a rate of 10°C/min and held at maximum tempera ture  for 30 min,  then the  furnace was  cooled inertially to  room temperature. The ceramic properties obtained are given  in Table 2.  RESULTS AND DISCUSSION  Starting powders  According to results obtained (see Table 1) for all pow ders measured r  p is less than the theoretical value. For ZrB2  powders  the  difference  is more marked  (for  example  r  p = 5.62 g/cm3 against r  t = 6,12 g/cm3), which may point to  48  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  TABLE 1. Starting Powder Main Properties  Material  Supplier, grade, TU or GOST  r  p, g/cm3  Phase composition  Particle size, mm  s  sp, cm2/g  Main impurities, wt.%  ZrB2  OOO KTM, TU 6-09-03-46-75  5.623  ZrO2 to 3 wt.%  1 - 40; dso = 9.2  2200  Al 0.11, Cu 0.17, Hf 1.0,  O 3.0, C 5.0  ZrB2  OOO Ékos-Ural, TU 6-09-03-46-75  5.820  ZrO2 to 7 wt.%  1 - 40; dso = 9.4  2800  Al 0.2, Cu 0.2, O 2.0, C 5.0  SiC  OOO ZAC, GSC, 64S, M5, GOST 26327  —  a-SiC  2 - 18; dso = 7.2  6675  C 0.27, Fe 0.09, O 0.65  SiC  AO VA, GSC, 64S, M5, GOST 26327  —  a-SiC  1 - 13; dso = 4.8  9898  C 0.04, Fe 0.2, O 0.5  Si3N4  OAO NEOMAT, TU 075-95 LR  —  Si3N4 amorph  0.05 - 0.10  >20 - 104  Fe 0.1, Al 0.1, O 4.0  TiSi2  DCRP (from 1988), TU 6-09-03-370-74  3.973  Tr. TiSi, Si  1 - 50  —  —  MoSi2 OOO EvroKhimInvest, TU 6-09-03-395-74  4.642  Tr. Mo5Si3  1 - 50  —  —  ZrSi2  OOO Ékos-Ural, TU 6-09-03-15-75  6.027  ZrO2 to 3 wt.%,  tr. ZrO, Si, SiO2  1 - 50  —  —  * GSC is green silicon carbide; ZAC is Zaporozh’e Abrasive Combine; VP is Volga Abrasive Plant; DCRP is Donetsk Chemical Reagent Plant;  Tr. is traces (<1 wt.%).  5  Ceramic machined surface was measured in a TR-200 profilo meter according to GOST 2789.  \\x0c', 'presence in powders of dense impurities, often x-ray-amor phous. Actual impurities for borides are B2O3 and boric acid,  and for  transition metal silicides it  is silicon dioxide. Com bined with data for elemental analysis the B2O3 content  in  ZrB2 powder is at a level from 3 to 6 wt.%.6 Presence of oxy gen  impurity with  quite  a  significant  amount  of  carbon  makes it possible to suggest  that  the ZrB2 powder was pre pared by carboor borothermal methods (oxide reduction).  Results of XPA showed that ZrB2 powder contains ZrO2  (baddeleyite)  in an amount  from 3 to 7 wt.% (see Table 1).  No other crystalline impurities were detected in ZrB2 pow der. The  calculated  crystal  lattice  parameter  of ZrB2  of  both  powders within  the  limits  of  error  (a = 3.167(1) Å,  c = 3.529(1) Å)  conform  well  with  published  data  (a = 3.168 Å,  c = 3.530 Å [1 - 3]). According  to  electron  microscope data the shape of original ZrB2 powder particles  is rounded, and the grain size is 1 - 8 mm. Particles were col lected into agglomerates of  indefinite shape with a size of  10 - 40 mm. These data are in good agreement with results of  laser diffraction for original powders (see Table 1).  Mixtures after milling in a planetary mill  After milling powders of a ZrB2-SiC mixture under the  conditions selected the average aprticle size decreased from  about 9 to about 4 mm and the limits for particle size distribu tion also shifted into a region of smaller sizes (from 0.1 to  10 mm). No signs of  forming particle  aggregates  (10 mm)  were revealed by results of laser diffraction.  Hot pressing  Rapid  shrinkage  under HP conditions  normally  com menced at 1300°C and was  entirely complete  at 1800°C,  which is connected with formation of phases with reduced  viscosity in the ZrO2-SiO2-B2O3 system. According to data  in [12 - 17] shrinkage normally commences at 1400°C. With  a synthesis temperature above 1900°C there is disilicide de composition and above 200°C visible inhomogeneity devel ops in ZrB2-SiC ceramic.  Ceramic after HP  According to hydrostatic weighing the density of  ce ramic was 97 - 99% (depending on composition), whereas  open porosity was from 0.1 to 3.0% (see Table 2). Minimum  porosity (<0.1%) was obtained in a specimen with molybde num disilicide  (No. 5a). Water  absorption of many speci mens was at the level of hundredths of a percent. It should be  noted that under similar HP conditions (~1800°C, holding for  30 min) ZrB2 powder produced by Ékos-Ural without addi tive sintered into a monolith with relative density more than  87%. A specimen is covered with a network of penetrating  cracks, which makes it impossible to test properties. Individ ual cracks in sintered ceramic were also observed in a pair of  ZrB2  (Ékos-Ural) and SiC (ZAK), which is connected with  the relatively high ZrO2 impurity content.  Results  of  optical microscopy  showed  that  isolated  grains of SiC are present within the matrix of ZrB2. Grain  size (irregular shape)  is about 6.0 mm for ZrB2 and 4.8 mm  for SiC, i.e., there is no marked increase in grain size for the  main phases in the course of the HP regime selected. Accord ing to electron microscopy data grains are visible in almost  all specimens identified as ZrO2 phase (with a volume fac tion of up to 4%). In the course of microsection preparation  spalling  proceeds  predominantly  through  silicon  carbide  grains. Results of XMRA showed that  impurities Al, Ca, Fe,  and Cr (up to 0,1 wt.%) are present. According to XPA data  all ceramic specimens consists of ZrB2 (hex.) and SiC (hex.  and cub.), and ZrO2  (mon.)  is observed in a small amount.  Presence of  any additional phases  apart  from the original  composition is extremely unlikely.  A key  characteristic  of  composite  ceramic  based  on  ZrB2-SiC is ultimate strength in bending at  room and ele Ultrahigh-Temperature Ceramic Based on ZrB2-SiC: Preparation and Main Properties  49  TABLE 2. Sintered Ceramic Specimen Main Properties  Specimen  number  Composition,  vol.%  r  a, g/cm3  r  a/r  t,  %  P,  %  w,  %  s  ben,  MPa  HV,  GPa  KIc, MPa·m1/2  a  20-1100, 10-6 K-1  l,  W/(m·K)  1  ZrB2-31SiCw-8Si3N4  4.92  99.0  3.1  0.63  525  10.6  3.5  6.4  83.2  2  ZrB2-9Si3N4  5.40  92.3  0.3  0.06  —  10.3  3.9  —  —  3  ZrB2-30SiC  5.13  98.0  0.32  0.06  635  14.6  5.4  6.2  90.2  4  ZrB2-29SiC-6TiSi2  5.09  98.8  0.28  0.03  400  13.0  5.8  —  —  5a  ZrB2-29SiC-5MoSi2  5.14  97.0  0.08  0.02  420  15.3  5.9  —  —  5b  ZrB2-29SiC-5MoSi2  5.08  96.0  0.32  0.06  535  14.5  4.8  6.3  53.8  6a  ZrB2-30SiC-5ZrSi2  4.49  86.5  11.93  2.65  320  7.3  5.1  —  —  6b  ZrB2-30SiC-5ZrSi2  5.03  97.0  0.29  0.05  590  14.3  5.5  —  —  6  In the course of treatment under vacuum in the initial HP stages  boron oxide partly evaporates. The quantitative  residue  in ce ramic was not evaluated in this work.  \\x0c', 'vated temperature; in the present work it was measured at 20,  1200, and 1400°C.7  A sharp drop in strength at 1400°C and existence of plas ticity for individual specimens is explained by an amorphous  phase at grain boundaries. It  is possible to retain strength at  these temperatures only by reducing the proportion of oxide  phases in original powders. According to [3, 18] glass phase  is the main sintering agent  in the course of HP. Similar be havior (a drop in ceramic strength >1200°C) is also observed  by authors overseas ([12 - 18], Fig. 1).  Data for LTEC (Fig. 2) are in good agreement with indi ces provided in [3, 12, 16, 17]. The value of LTEC is hardly  sensitive to a change in UHTCM composition within the lim its up to 5 vol.%. Within the limits of measurement error the  temperature  paths  of  thermal  expansion  curves  agree  for  specimens Nos. 1, 3, and 5b. Indices obtained are close to the  LTEC for  ceramic  based  on  aluminum oxide  (a  20-900 =  5 ´ 10-6 - 9 ´ 10-6 deg-1), and at  the same time they are ap proximately half those for classis heat-resistant metal alloys  such  as  Inconel  718  (a  20-600 = 14.9 ´ 10-6 deg-1). For  ce ramic based on ZrB2-SiC at above 1300°C an increase in  LTEC almost ceases [3, 16].  Heat capacity, measured in specimens Nos. 1, 3, and 5b,  increased sharply from 0.45 - 0.47 to 0.68 - 0.73 J/(g·deg) at  25 and 700°C respectively. This is within the limits of 3% er ror  for  the method and agrees with published data for  the  composition ZrB2-30% SiC [16]. Thermal conductivity of  this group of specimens has marked scatter in values depend ing on the raw material used. For specimens Nos. 1 an 3 it  decreases sharply from 34 - 39 to 16 - 18 mm2/sec at 25 and  700°C respectively, whereas  for  a  specimen with MoSi2  (No. 5b)  it  decreases more  smoothly,  i.e.,  from 22 to  13 mm2/sec in the same temperature range. Thermal conduc tivity at high temperature emerges into a plateau with a value  of about 15 mm2/sec for composition ZrB2-30% SiC [16].  For specimen No. 1 the calculated thermal conductivity de creases uniformly from 83 to 65 W/(m·deg) with an increase  in temperature from 20 to 1100°C respectively (Fig. 3).  It  50  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  7  The temperature of 1400°C for UHTCM is its “boundary temper ature”. Above 1400°C there is melting of silicon and marked soft ening of the glass phase of the composition Zr-B-Si-O at grain  boundaries. Achievement of high strength at these temperatures is  technically impossible [3], and is currently being studied. For use  in hypersonic aircraft  it  is necessary that UHTCM have an ulti mate  strength  in  bending  not  less  than  200 MPa  at  above  1500°C [3].  Fig. 1. Ultimate strength in static bending of ceramic specimens  based on ZrB2-SiC with different additions: unfilled symbols are  data of the article authors (one symbol — one test); filled symbols  are published data (for convenience some filled symbols at  room  temperature are moved the left).  Fig. 2. Average LTEC for ceramic specimens Nos. 1, 3, and 5b (un filled symbols) compared with published data (filled symbols).  Fig. 3. Dependence of thermal conductivity of ceramic specimens  based on ZrB2-SiC with different additions (Nos. 1, 3, 5b) on tem perature (unfilled symbols) compared with published data (——,  -).  \\x0c', 'should be noted that by comparing the temperature depend ence with data provided in [3, 16] values obtained in the  present work for  the  composition with addition of Si3N4  (No. 1) are somewhat higher than expected. Relatively high  thermal conductivity of specimen No. 1 is explained by three  factors. First,  its high relative density. Second, a relatively  low zirconium dioxide  impurity,  and  also  use within  the  composition  of  specimens  No. 1  of  silicon  carbide  microfiber. For specimen No. 5b the thermal conductivity is  somewhat  lower, which is  connected with an increase  in  ZrO2 impurity is starting raw material.  Microhardness  of  a  specimen with  addition  of  SiO2  (No. 5a) is comparable with that of hard alloy VK8 WC-8Co  (Fig. 4). For all specimens of ceramic with added disilicides  KIc > 4.8 MPa·m1/2 (even for specimen No. 6a with a density  of not more than 86%), which is comparable with data ob tained previously for ZrB2-SiC ceramic [7] sintered by the  SPS method. The value of K1c obtained by the authors of the  present article is markedly higher  than that encountered in  publications of  the value for ZrB2-SiC ceramic. For exam ple,  in  publications  [12, 18]  values  are  shown  of  4.3 - 4.6 MPa·m1/2. Data for K1c given in different work are  difficult  to compare with each other due to existence of sev eral  independent measurement methods and calculation of  ceramic material crack resistance. It should be noted that the  size of grains in ceramic is smaller by a factor of 5 - 10 than  the impression of an indenter with a load of 1 - 3 kg,  i.e.,  only tests with increased loads (³9.8 N) are indicative.  The weight  increase  for UHTCM is  not more  than  0.3 wt.% up to 1500°C in a stream of commercially pure ni trogen. On the DSC curve  there were no features which  points to absence of phase transitions in specimens. Holding  in air at up to 1000°C inclusively does not change the exter nal appearance of specimens and their weight.  In the range  1100 to 1300°C for all specimens there is a continuous matt  film (consisting of borosilicate glass);  the weight  increase in  not more than 0.3% (~2 mg/cm2).  At  1400 - 1500°C in  specimens  there  are  individual  inhomogeneities and swelling of oxide film. The weight  in crease for specimens Nos. 1, 3, 4, 5a, 5b, and 6b remains at  the level of 0.2 - 0.3% specimen No. 6a (with addition of  ZrSi2 and density ~ 86%) has a markedly greater weight  in crease,  i.e., up to 1%. At 1600°C weight  increase for speci men No. 4 (with added TiSi2) was 2.6%, for specimen No. 5a  (with addition of MoSi2)  it  increased up to 0.5%, and for  specimen No. 6a up to 1.6%. Specimens Nos. 4, 5a, and 6a  before and after oxidation at 1700°C are shown in Fig. 5.  Only for a specimen with addition of MoSi2 is a continuous  transparent oxide  film retained;  for  the  rest of  specimens  there is formation of an oxide film white in color and its sep aration. At  the  surface of  specimens Nos. 4 and 6a  local  cracks and blisters arise, which is apparently connected with  a  reduced viscosity of  the phases  formed of  composition  Zr-Ti-Si-O-B. Thus,  addition  of Si3N4  and TiSi2  in  an  amount of 5 vol.% to the UHTCM composition reduces oxi dation resistance considerably at above 1700°C. Addition of  MoSi2  in contrast  improves oxidation resistance compared  with the basic composition No. 3 (ZrB2 - 30% SiC). The re sult obtained a at qualitative level agrees with data obtained  previously [3, 12 - 18].  In comparing the main HP indices for HP ceramic with  the properties of ceramic previously prepared by us using the  SPS method [7] it is seen that that due to an increase in rela tive density from 92 - 93 to 97 - 99% it was possible to raise  microhardness and KIc considerably. Transfer  to finer SiC  powder  (with d50  from 7.2 to 4.8 mm) made it possible to  achieve the best properties with the same nominal composi tion. A further increase in indices was facilitated with addi tion MoSi2 and ZrSi2 in amount of 5 vol.%. Use of very fine  silicon nitride powder as a sintering agent appeared to be in valid. For ceramic quite low hardness and K1c (see Table 2).  It may be suggested that  improvement of chemical and  phase purity of starting powders of ZrB2 and SiC, a further  reduction in SiC grain size (to 1 mm and less), and optimiza tion of the HP parameters (in order to achieve r  a/r  t > 99.5%)  Ultrahigh-Temperature Ceramic Based on ZrB2-SiC: Preparation and Main Properties  51  Fig. 4. Vickers microhardness and K1c for ceramic specimens based  on ZrB2-SiC with different additions (Nos. 3, 6a, 4, 5a) compared  with published data.  Fig. 5. Specimens of ceramic based on ZrB2-SiC with additions  (left  to right) TiSi2 (No. 4), MoSi2 (No. 5a) and ZrSi2 (No. 6a) be fore (a) and after (b) oxidation at 1700°C with exposure at the maxi mum temperature for 30 min.  \\x0c', 'will make it possible to increase ceramic strength at moder ate and elevated temperatures.  CONCLUSION  By using entirely domestic raw material  it  is possible to  prepare dense (>97%) ceramic based on ZrB2-SiC by an HP  method. The best properties are shown by the composition  ZrB2 - 29% SiC - 5% MoSi2. This ceramic has high hard ness (~15 GPa) and K1c (up to 5.9 MPa·m1/2). On holding in  air  at  1800°C its weight  increase was  less  than  1 wt.%  (<10 mg/cm2). However, with  an  increase  in  temperature  from room to 1400°C there is a sharp reduction in strength in  three-point bending,  i.e., from 500 to 200 MPa. This occurs  due to presence within ZrB2 powder of  impurities  (mainly  oxygen).  The authors of the article thank R. A. Mironov, T. A. Mish nova, T. S. Frolova,  I. M. Rudykin, and G. M. Bagreeva for  participation in the work.  REFERENCES  1. G. V. Samsonov, T. I. Serebryakova, and V. A. Neronov, Borides  [in Russian], Atomizdat, Moscow (1975).  2. T. I. Serebryakova, V. A. Neronov, and P. D. Peshev, High-Tem perature Borides [in Russian], Metallurigiya, Moscow (1991).  3. W. G. Fahrenholtz, E. J. Wuchina, W. E. Lee, and Y. Zhou (edi tors), Ultra-High Temperature Ceramics. Materials for Extreme  Environment Applications, Wiley, New Jersey (2014).  4.  . K. Sonber and A. K. Suri, “Synthesis and consolidation of zir conium diboride:  Review,”  J.  Adv.  App. Ceram.,  110(6),  321 - 324 (2011).  5. E. P. Simonenko, D. V. Sevast’yanov, N. P. Simonenko, et al.,  “Simonenko, E. P. Promising ultra-high temperature ceramic  materials  for aerospace applications,” Russ. J.  Inorg. Chem.,  58(14), 1669 - 1693 (2013).  6. R. A. Andrievskii, “Nanostructured diborides of titanium, zirco nium, and hafnium: synthesis, properties, dimensional effects,  and stability,” Uspekhi Khimii, 84(5), 540 - 554 (2015).  7. L. A. Chevykalova,  I. Yu. Kelina,  I. L. Mikhal’chik,  et  al.,  “Preparation of ultra-high temperature ceramic material based  on zirconium boride by SPS method,” Refract. Indust. Ceram.,  54(6), 455 - 462 (2014).  8. D. V. Grashchenkov, O. Yu. Sorokin, Yu. E. Lebedeva,  and  M. L. Vaganova, “Specific features of sintering of HfB2-based  refractory ceramic by hybrid spark plasma sintering,” Russ. J.  App. Chem., 88(3), 386 - 393 (2015).  9. S.  S. Ordan’yan, V.  I. Rumyantsev, D. D. Nesmelov,  and  D. V. Korablev, “Physicochemical basis of creating new ceram ics with  participation  of  boron-containing  refractory  com pounds  and  its  practical  implementation,”  Refract.  Indust.  Ceram., 53(2), 108 - 111 (2012).  10. V. A. Sokolov, “Production condition and raw material base of  zirconia  refractories  in Russia,” Novye Ogneupory, No. 11,  13 - 17 (2013).  11. G. S. Khodakov, “Method for measuring specific surface of ver ified powders by gas filtration,” Kolloid. Zh., 57(2), 280 - 282  (1995).  12.  J. F. Justin and A. Jankowiak, “Ultra-high temperature ceram ics: densification, properties  and thermal  stability,”  J. Aero space Lab., No. 3 (2011).  13. N. Gupta, A. Mukhopadhyay, K.  Pavani,  and  B.  Basu,  “Spark-plasma sintering of novel ZrB2-SiC - TiSi2 composites  with better mechanical properties,” Mater. Sci. Eng. A., 534,  111 - 118 (2012).  14. S.-Q. Guo, T. Nishimura, T. Mizuguchi, and Y. Kagawa, “Me chanical properties of hot-pressed ZrB2-MoSi2 - SiC compos ite,” J. Europ. Ceram. Soc., 28, 1891 - 1898 (2008).  15. O. N. Grigoriev, B. A. Galanov, V. A. Kotenko, et al., “Mechan ical properties of ZrB2-SiC (ZrSi2) ceramics,” J. Europ. Ceram.  Soc., 30, 2397 - 2405 (2010).  16.  J. W. Zimmermann, G. E. Hilmas, W. G. Fahrenholtz, et al.,  “Thermophysical properties of ZrB2 and ZrB2-SiC ceramics,”  J. Amer. Ceram. Soc., 91, 1405 - 1411 (2008).  17. F. Monteverde, S. Guicciardi,  and A. Bellosi,  “Advances  in  microstructure and mechanical properties of zirconium diboride  based ceramics,” Mater. Sci. Eng. A, 346, 310 - 319 (2003).  18. R. Loehman, E. Corral, H. P. Dumm, P. Kotula, and R. Tandon,  “Ultra-high temperature ceramics for hypersonic vehicle appli cations,” Sandia Report, USA (2006).  52  P. S. Sokolov, A. V. Arakcheev, I. L. Mikhal’chik, et al.  \\x0c']"
},{
  "_id": 282,
  "PDF": "Understanding the oxidation behavior of Ta–Hf–C ternary ceramics at high temperature.pdf",
  "Text": "['Journal Pre-proof  Understanding the oxidation behavior of Ta-Hf-C ternary ceramics at high temperature  Jian Zhang, Song Wang, Wei Li, Yiping Yu, Jinming Jiang  PII:  DOI:  S0010-938X(19)31349-6  https://doi.org/10.1016/j.corsci.2019.108348  Reference:  CS 108348  To appear in:  Corrosion Science  Received Date:  1 July 2019  Revised Date:  26 October 2019  Accepted Date:  15 November 2019  Please cite this ar ticle as: Zhang J, Wang S, Li W, Yu Y, Jiang J, Understanding the oxidation behavior of Ta-Hf-C ternary ceramics at high temperature, Corrosion Science (2019), doi: https://doi.org/10.1016/j.corsci.2019.108348  This is a PDF file of an ar ticle that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the ar ticle. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal per tain.  © 2019 Published by Elsevier.  \\x0c', 'Understanding the oxidation behavior of Ta-Hf-C ternary ceramics at high temperature   Jian Zhanga, Song Wanga, *, Wei Lia, Yiping Yua, Jinming Jiangb   a Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University   of Defense Technology, College of Aerospace Science and Engineering, Changsha 410073, China   b Department of Basic Sciences, Air Force Engineering University, X ian 710000, China   *Corresponding author. E-mail: wangs_0731@163.com   Journal Pre-proof  rate kp is ranked in the order of 1TaC-3HfC ＜ 1TaC-1HfC ＜ 1TaC-4HfC ＜ 4TaC-1HfC.   The new three-dimensional Hf6Ta2O17-Ta2O5 hyper-eutectic scale and the formation of a dense pure   Hf6Ta2O17 inner layer combined with a continuous amorphous Hf-Ta-O-C transition layer are found   The oxidation process of Ta-Hf-C ternary ceramics exhibit composition dependence: the oxidation   Highlights   \\uf06c   \\uf06c   \\uf06c   The study of   the oxidation   resistance of Ta-Hf-C ceramics is significant for their   application in real engineered systems, such as thermal structural components of a   hypersonic vehicle, etc.   to be responsible for the improved oxidation performance exhibited in 1TaC-1HfC and 1TaC-3HfC,   respectively.   1             \\x0c', '\\uf06c   We believe that Ta-Hf-C ternary ceramics will have better comprehensive properties when the   TaC/HfC molar ratio is between 1/1 and 1/3.   Abstract:   1TaC-3HfC exhibit considerably improved oxidation resistance compared to TaC and HfC individual   carbides along with 4TaC-1HfC and 1TaC-4HfC ternary ceramics. The results obtained herein should   Ta-Hf-C ternary ceramics are studied systematically at 1400-1600°C in air under isothermal conditions   for the first time. The oxidation process is demonstrated to be composition dependent: 1TaC-1HfC and   important   life-limiting factor. In   this work,   the oxidation mechanisms of various compositions of   pave the way to design advanced oxidation-resistant Ta-Hf-C ternary ceramics for high-temperature   For high-temperature applications, poor oxidation performance of Ta-Hf-C ternary ceramics is an   Journal Pre-proof  Currently, the keen interest in hypersonic aircraft is driving fierce worldwide competition to develop   strength, and good oxidation resistance [1-3]. However, only few known materials possess the essential   applications.   1. Introduction   Keywords: Ta-Hf-C ternary ceramics; Oxidation; Microstructure; Kinetics; Mechanism   next-generation structural materials that have high-temperature resistance (>3000°C), high mechanical   refractoriness to withstand such ultra-high temperatures, let alone the simultaneous corrosive action of an   oxidizing environment. As solid solution of tantalum carbide (TaC) and hafnium carbide (HfC), Ta-Hf-C   ternary ceramics are new possibilities in this field of research because they have (i) the highest melting   2           \\x0c', 'point   (almo st 4000°C   [4]) of all known materials,   (ii) excellent mechanical properties, and   (iii)   appropriate thermo-physical properties [5, 6].   However, some major limitations remain to be addressed in the application of Ta-Hf-C systems.   Ta-Hf-C ternary ceramics are not often obtained as single-phase materials because their strong covalent   bonds and low self-diffusion coefficients make sintering of these ceramics particularly challenging. Some   solid-solution process and samples of high bulk density remains difficult [7, 8]. Effective approaches to   mitigate the solid-solution issue and improve the density of Ta-Hf-C ternary ceramics envisage the   studies have considered the sintering of initial TaC/HfC mixable powders; however, achieving sufficient   preparation of nano-scale Ta-Hf-C solid-solution powders followed by sintering. Based on this, in   Journal Pre-proof  Having solved the problems of preparing Ta-Hf-C ternary ceramics, another issue that we must   resistance properties of   these materials. Oxidation   resistance   the materials can survive under oxidizing conditions at high temperatures [12, 13]. However, to the best   of our knowledge, the literature on the oxidation behavior of Ta-Hf-C ternary ceramics is very scarce.   Previous studies indicate that the oxidation resistance of TaC and HfC individual carbides is no t satisfied   performance is most challenging in the development of carbides because it determines directly whether   previous work, we reported a novel joint process of solvothermal and hot pressing for preparing highly   dense single-phase Ta-Hf-C ternary ceramics [9, 10]. The obtained Ta-Hf-C samples were shown to   have a highly uniform crystal structure and excellent mechanical properties, thereby making such samples   potentially applicable in future engineering fields [11].   address   is evaluating   the oxidation   because no continuous protective oxide layer forms during oxidation at high temperature. TaC is oxidized   fully at temperatures as low as 850°C, and the oxide Ta2O5 is unstable, liquefying at high temperature   3     \\x0c', 'because of its low melting temperature (~1800°C) [14]. HfC has been shown to oxidize and form a   porous HfO2 scale at temperatures starting from 400°C [15]. It is only at very high temperatures (higher   than even 1600°C) that the HfO2 oxide scale can sinter quickly enough to provide reasonable oxidation   protection [16]. In this case, by combining TaC and HfC individual carbides, Ionescu et al. showed that   incorporating Ta   into HfC enhances   the   sintering ability of   the oxide   scale,   thereby   improving   oxidation resistance of all the compositions. However, none of these limited studies provided a clear   temperature of approximately 2800°C [18]. The solid-solution samples were found to have considerably   improved oxidation resistance compared to the pure carbide samples, and HfC-50vol% TaC had the best   and TaC) affect the oxidation behavior of Ta-Hf-C ternary ceramics by exposure to a plasma torch with a   considerably the oxidation resistance of (Hf1−xTax)C/SiC composites [17]. In addition, Zhang et al.   investigated how different compositions (HfC, HfC-20vol% TaC, HfC-50vol% TaC, HfC-80vol% TaC,   verdict on   ternary ceramics during   the oxidation kinetics and microstructural changes of Ta-Hf-C   Journal Pre-proof  In the present study, various compositions of Ta-Hf-C ternary ceramics (4TaC-1HfC, 1TaC-1HfC,   high-temperature oxidation; consequently, the oxidation mechanisms of various Ta-Hf-C compositions   are yet to be fully understood. Moreover, the important 1TaC-3HfC (a one-to-three molar ratio of TaC to   HfC) composition is yet to be investigated despite the belief that it has the best oxidation resistance   among Ta-Hf-C systems according to the oxidation results for Ta-Hf alloys [19, 20].   1TaC-3HfC, and 1TaC-4HfC) were oxidized at 1400-1600°C in air. The oxidation kinetics and the   microstructures of the oxide scales of these Ta-Hf-C ternary ceramics were investigated systematically   for the first time. The obtained detailed knowledge about the thermodynamic data and microstructure   evolution during high-temperature oxidation is expected to help obtain a good understanding of the   4     \\x0c', 'degradation of Ta-Hf-C ternary ceramics. In particular, detailed microstructural analysis by transmission   electron microscopy (TEM) revealed characteristics at the nano-scale that could not be detected by other   normally used characterization   techniques such as scanning electron microscope (SEM) and X-ray   diffraction (XRD). Several interesting discoveries that have never been reported previously are discussed   herein.   powders were produced by the joint processes of solvothermal treatment and carbothermal reduction   method, which was established in our previous work [9, 10]. The generated powders with average   2. Experimental   2.1. Ta-Hf-C ternary ceramics fabrication   Various compositions of nano-scale Ta-Hf-C (4TaC-1HfC, 1TaC-1HfC, 1TaC-3HfC and 1TaC-4HfC)   Journal Pre-proof  Small batches of the Ta-Hf-C powders (12-15 g) were place in graphite dies (20 mm diameter) lined   The resulting specimens were ground and cleaned, subsequently, several cylinders of size φ7mm ×   with graphite foil, and densified using hot pressed sintering (HP; Shanghai Chenhua instrument Co., Ltd.,   China). The HP conditions were: 2100 °C peak temperature, 70 MPa pressure, 30 min hold duration, and   3mm were cut from the discs for oxidation studies of the as-sintered Ta-Hf-C ternary ceramics. The   cylinder samples were placed at a certain angle of inclination in a zirconia crucible against its wall such   that only two edges of the specimens were in contact with the crucible. Isothermal oxidation tests of the   5   crystallite size of ~80nm, average particle (aggregate) size of ~400nm and oxygen content of 0.8 -1.5   wt.%.   free cooling followed.   2.2. Oxidation tests     \\x0c', 'samples were performed at 1400-1600 °C in air flowing by natural convection in a muffle furnace. After   the furnace was heated up to the setting temperatures, the samples were directly put into the furnace,   maintained at this temperature for 5min to 300min, and then they were taken out of the furnace to record   the mass. The oxidation fraction (α) of the samples were calculated by dividing the measured weight   increase by the theoretical one, as the following equation   (1)   where m0, mt and m∞ are   the weight of   the specimens before oxidation, after oxidation and   completely oxidized respectively. And the m∞ was calculated by assuming the complete conversion   of Ta-Hf-C to HfO2 and Ta2O5.   2.3. Characterizations   𝑚𝑡 − 𝑚0  𝑚∞ − 𝑚0  𝛼 =  Bulk densities of the sintered Ta-Hf-C ceramics were measured using the Archimedes principle with   Journal Pre-proof  The micro-morphologies before and after oxidation were analyzed on surface and polished cross   Crystalline phases of the oxidized ceramics were conducted using X-ray diffraction with a Bruker   D8 Advance using CuKα radiation (λ= 1.5406) operated at 30 mA and 40 kV (Saarbrucken, Germany).   distilled water as the immersion fluid. Pure component densities of TaC and HfC are defined to 14.50 and   12.70 g/cm3, respectively. And the theoretical densities of the Ta-Hf-C solid solution ceramics were   calculated using the rule of mixtures and is 14.14 g/cm3 for 4TaC-1HfC, 13.60 g/cm3 for 1TaC-1HfC,   13.15 g/cm3 for 1TaC-3HfC, 13.06 g/cm3 for 1TaC-4HfC. The calculated theoretical densities were used   in calculating the relative densities throughout this study.   sections by scanning electron microscopes(SEM, S4800, Hitachi Limited, Japan). Compositions of the   samples were obtained in the SEM using energy dispersive spectroscopy (EDS).   6         \\x0c', 'A transmission electron microscopy (TEM, JEM-2100, Japan Electronics Co., Ltd, Nagano, Japan),   operated at 200keV accelerating voltage and also equipped with EDS, was used to characterize the   cross-sections of oxidized 1TaC-1HfC and 1TaC-3HfC. TEM specimens were extracted from spec ific   locations by in situ lift-out of foils nano-machined using a focused ion beam microscope (FIB; Helios   NanoLab 600i, FEI, Hillsboro, OR).   Ta-Hf-C ternary ceramics (4TaC-1HfC, 1TaC-1HfC, 1TaC-3HfC, and 1TaC-4HfC) are presented in Fig.   the basic information in this section for the sake of clarity. The XRD patterns of the four compositions of   1(a). Mutual diffusion of TaC and HfC occurred in each sample, and thus five typical diffraction peaks of   the cubic phase are observed in isolation. The peaks move to lower 2θ values as the HfC content is   3. Results   3.1. Basic information about Ta-Hf-C ternary ceramics   We have characterized the obtained Ta-Hf-C ternary ceramics in previous work [11], but we present   Journal Pre-proof  Table 1 show that the obtained Ta-Hf-C ternary ceramics are near full densification, with some   evident, and Ta, Hf, and C are distributed homogeneously throughout the representative 1TaC-1HfC   sample, Fig. 1(c). These results indicate that a uniform and sufficient solid-solution process was achieved,   increased, indicating the formation of solid solution, Fig. 1(b). A high-density surface microstructure is   and the obtained samples are characterized as being single phase and highly dense.   improvement upon increasing the TaC content. Furthermore, they have excellent mechanical properties,   particularly hardness values of 30-35 GPa, making them some of the hardest ultra-high-temperature   ceramics. Furthermore, their thermal conductivities are relatively high, with each sample exceeding 23   W·m−1K−1. As for the preparation method and additional details about the samples, these are described in   7     \\x0c', 'for the comparison samples of pure TaC and HfC. The α values for 1TaC-1HfC and 1TaC-3HfC are   Comparatively, the α values for the 4TaC-1HfC and 1TaC-4HfC compositions are only slightly below   our previous reports [9-11].   3.2. Oxidation behavior   The oxidation behavior at different temperatures has been studied in terms of the variation in   oxidation fraction (α) with time (t). Fig. 2(a) shows the values of α measured for the four compositions of   Ta-Hf-C ternary ceramic after 30 min of exposure at 1400°C, 1500°C, and 1600°C, together with those   those for pure HfC, indicating that excessively high or low TaC content in Ta-Hf-C systems can only   slightly improve the HfC oxidation resistance. In contrast, comparing the α values for the Ta-Hf-C   clearly much lower than those for the other four samples, thereby indicating superior oxidation resistance.   varying degrees by the formation of Ta-Hf-C ternary ceramics even with different compositions.   Journal Pre-proof  the plots of α as a function of   Fig. 2(b)-(e) shows   temperatures of 1400-1600°C in air for up to 300 min. During the initial oxidation stage at each   temperature, the α value for each sample increases rapidly, and the oxide scale forms quickly on the   time increases at each temperature, the α value of each specimen increases slowly, indicating an oxide   passivation stage. Meanwhile, for a given oxidation temperature and time, the α values of the samples   decrease gradually as the HfC content increases from 4TaC-1HfC to 1TaC-3HfC, although they then   increase considerably for 1TaC-4HfC, implying that the 1TaC-3HfC composition has the best oxidation   resistance herein.   8   ternary ceramics with those of pure TaC shows that the oxidation resistance of pure TaC is improved in   t for   the Ta-Hf-C   ternary ceramics at   substrate surface, thereby preventing further oxidation to a certain extent. Subsequently, as the oxidation     \\x0c', '3.3. Microstructure upon oxidation   To   investigate   the phase compositions of   the oxidation products, Fig. 3 shows XRD patterns   obtained from the oxide scales of all the investigated samples after oxidization at 1500°C in air for 30 min.   The oxidized 4TaC-1HfC ceramic exhibits the formation of tri-Ta2O5 and traces of o-Hf6Ta2O17 upon   oxidation, Fig. 3(a). However, no HfO2 is detected, indicating that all the Hf in the oxide layer existed in   changes in the content of these two oxides. As for the 1TaC-3HfC composition, the peaks of tri-Ta2O5   2.5-3.5, and   this ratio   is   treated as the phase boundary for forming Hf-Ta-O solid solution [21].   Consequently,   is exactly right   for forming a   the form of o-Hf6Ta2O17 solid solution oxide. Upon increasing the HfC content, the phase compositions of   the 1TaC-1HfC oxide scale are the same as that of 4TaC-1HfC, as shown in Fig. 3(b). The peak intensity   disappear and only those of o-Hf6Ta2O17 can be detected, implying that the oxide scale comprises pure   o-Hf6Ta2O17. According to Spridonov et al., the Hf/Ta ratio in single -phase Hf-Ta-O solid solution is   of tri-Ta2O5 decreases while that of o-Hf6Ta2O17 increases, suggesting that there are corresponding   Journal Pre-proof  Figs. 4-10 show SEM and TEM   the microstructure as well as   the value of Hf/Ta=3/1 for   the 1TaC-3HfC sample   images of   the corresponding   single-phase o-Hf6Ta2O17 structure in the 1TaC-3HfC composition. In this case, for 1TaC-4HfC with   excess HfC, the oxide scale is found to comprise o-Hf6Ta2O17 and m-HfO2 in the corresponding XRD   30 min (owing to the incapacity of EDS for light elements, the content of oxygen was not analyzed) .   Clearly, the microstructure changes considerably with the Hf/Ta molar ratio.   4TaC-1HfC    The sample surface comprises a continuous molten liquid layer after oxidation, Fig.   9   pattern.   energy-dispersive spectroscopy (EDS) results for the Ta-Hf-C ternary ceramics oxidized at 1500°C for       \\x0c', '4(a), on top of which a small amount of isolated micron-sized crystals, indicating low wettability, is   commonly observed, Fig. 4(b). In addition, some uniformly distributed pores are detected, which are   caused by the release of gas byproducts generated during oxidation. Combined with EDS and XRD   analysis, Fig. 4(c), the molten liquid (point 1) is mainly identified as Ta2O5 and the isolated crystals (point   2) as Hf6Ta2O17.   oxidation scale is approximately 200 μm thick. The enlarged image of the oxide layer shows that it   where V is the molar volume, M is the molar weight, and ρ is the theoretical density. The larger the   scale broke away from the substrate, indicating that the oxidation layer bonds weakly with the carbide   matrix, Fig. 4(f). According to the Pilling-Bedworth theory [22], when an oxide forms at a carbide/oxide   comprises mainly compact Ta2O5 glass with small bright Hf6Ta2O17 precipitates and dark micro-pores   distributed therein Fig. 4(e), which is coincident with the analysis of the surface morphology. The oxide   An overview of the polished cross section reveals a mono-layered structure, Fig. 4(d), and the   Journal Pre-proof  𝑛 ∙ 𝑀𝑀𝑒𝐶 ∙ 𝜌𝑀𝑒𝑂  expressed with   the   𝑀𝑀𝑒𝑂 ∙ 𝜌𝑀𝑒𝐶  change due   𝑅𝑃𝐵 =  𝑉𝑀𝑒𝑂  to   the   formation of   the oxide   can be   =  𝑉𝑀𝑒𝐶  (2)   difference of RPB from 1, the higher the interfacial stress exists in the scale and more likely the oxide scale   degrades its oxidation resistance ability. Based on Eq. (2), the value of RPB for Ta2O5 and Hf6Ta2O17 from   interface,   the volume   Pilling-Bedworth ratio (RPB):   the oxidation of 4TaC-1HfC is calculated as 1.91, which is so large that volume expansion of the oxide   layer leads to high interfacial stress, thereby causing the oxide layer to peel off.   1TaC-1HfC    The oxidation surface possesses a dense matrix, again with Ta2O5 molten liquid   (point 3) as a continuous phase, but with considerably more Hf6Ta2O17 crystals (point 4), Fig. 5(a)-(c). By   10         \\x0c', 'contrast with the 4TaC-1HfC sample, few considerable defects can be observed in the surface scale,   indicating that the micro-pores can be infilled by Ta2O5 liquid, sealing defects while allowing the escape   of gaseous product.   The polished cross-section gives further indications, Fig. 5(d)-(f). The thickness of the oxide layer is   reduced to around 180 μm, Fig. 5(d), suggesting that the 1TaC-1HfC composition has a lower oxidation   Fig. 5(f).   rate than that of 4TaC-1HfC, which agrees well with the results for the oxidation mass gain. The   dispersed Hf6Ta2O17 precipitates do not exist in isolation anymore. Instead, a network of Hf6Ta2O17   precipitates and Ta2O5 molten liquid interpenetrate each other in the oxide scale , presenting a typical   “porous-framework/liquid-infill” structure, Fig. 5(e). Additionally, good interfacial bonding of the oxide   layer to the substrate is observed, which can be attributed to the substantial reduction in RPB (RPB=1.70),   To investigate further the microstructures of the highly dense oxide scale of 1TaC-1HfC, a focused   Journal Pre-proof  As a rule of thumb, bright regions correspond to Hf6Ta2O17 oxides and dark ones correspond to Ta2O5   glass. The Hf6Ta2O17 and Ta2O5 phases (marked as 1 and 2, respectively) present a great variation of   composition with complementary decreased or increased content of the metallic elements Ta and Hf: if   ion beam (FIB)/lift-out technique was used to prepare a site -specific thin sample, whereupon TEM   analysis was conducted. Fig. 6(a) shows the scanning transmission electron microscope (STEM) image   and the associated EDS element mapping results for Ta, Hf, and O of the site-specific thin sample. The   fine structure of the “porous-framework/liquid-infill” morphology is revealed clearly in the STEM image.   one is low, then the other is high. Meanwhile, the distributions of Hf, Ta, and O in the Hf6Ta2O17 and   those of Ta and O in the Ta2O5 are almost random with no segregation, indicating that both phases are   11     \\x0c', 'uniform single phase. Furthermore, high-resolution transmission electron microscope (HRTEM) and   corresponding selected area electron diffraction   (SAED) analysis show   the crystal structures and   interfacial combination (marked as 3) of Hf6Ta2O17 and Ta2O5. The Hf6Ta2O17 (Ta2O5) is verified as the   orthorhombic (triclinic) phase with lattice parameters a=4.941 Å, b=86.71 Å, and c=5.264 Å (a=3.801 Å,   b=3.785 Å, and c=35.74 Å), being consistent with the results of the XRD patterns, Fig. 6(b) and (c).   Furthermore, the change from one phase to the other occurs over one (or a few) interatomic distance only,   phases detected,   showing   thereby showing strongly   that   structure.   1TaC-3HfC   Above all, the phase interface between the Hf6Ta2O17 and Ta2O5 is clean and planar with no intermediary   (e).   layer   is exactly a eutectic   sharp morphology with good   interfacial matching, Fig. 6(d) and   the porous-framework/liquid-infill oxide    Unlike the two aforementioned compositions (4TaC-1HfC and 1TaC-1HfC) that   Journal Pre-proof  Moving inward from the front surface, a three-layer structure with a fully oxidized thickness of   approximately 130 μm is observed, Fig. 7(d)-(f). The outermost Hf6Ta2O17 layer comprises a porous outer   one and a dense inner one, Fig. 7(e). There are no visible impenetrating cracks in this scale, indicating   have high Ta content, the oxide layer of 1TaC-3HfC comprises isometric crystals of a single morphology,   with no apparent mixed oxides coexisting, Fig. 7(a). Although several cracks are obvious in the surface   layer, a high-densification sintered morphology can be observed   in   the uncracked area, Fig. 7(b).   Furthermore, the XRD and EDS results, Fig. 7(c), show that the phase composition of this layer can be   defined as single-phase Hf6Ta2O17 ternary oxide.   that the Hf6Ta2O17 layer slows down the diffusion of oxygen to some degree. The second layer, namely   the 2-μm-thick transition layer, has a dense structure in which there are no cracks or voids. Furthermore,   12       \\x0c', 'transition layer of 1TaC-3HfC is also an oxygenand carboncontaining material and is the key factor in   the excellent oxidation. Nevertheless,   to   the precise character of   the   between   the   oxycarbide transition layer reported in previous paper remains ambiguous ; for example, the form and   amount of C incorporated in the oxide-based matrix are unknown, and the crystal structures of the   it appears to adhere well to the residual carbide and the outer oxide , there being no cracking, spalling, or   separation at any of its interfaces with either of the other two layers , as shown in Fig. 7(f). Additionally,   the good binding can be attributed to a further reduction in RPB (RPB = 1.38). Finally, the third (i.e.,   innermost) layer is the unaffected substrate.   Previous work involving oxidation analysis of HfC individual carbide has shown that the HfCxOy   transition layer is a very important barrier to oxygen diffusion [16, 23]. From that, we infer that the dense   the best of our knowledge,   Journal Pre-proof  the C concentration   layer. Meanwhile,   the unaffected   transition   layer and   increases   element analysis results, we define the average empirical formula of Hf0.70Ta0.23O0.85C for the Ta-Hf-O-C   continually in the oxide layer and the transition layer and then remains constant in the unaffected layer.   This shows strongly that C and O coexist in the transition layer. Furthermore, based on the quantitative   oxide-based matrix is under-reported. To compensate for these limitations, we used wavelength-dispersive   spectroscopy (WDS), Raman spectroscopy, and TEM for additional characterization studies. WDS   line-scan analyses revealed gradients in the C and O contents across the transition layer, Fig. 8(a). The O   concentration decreases steeply at the interface between the oxide layer and the transition layer and at that   oxycarbide. The Raman spectra for the transition layer show two active modes, namely the D band and G   band, Fig. 8(b). The ID/IG ratio is found to be related to the stages of disorder and amorphization of C   materials. In the present transition layer, ID/IG is only 0.75, indicating that the C in Hf0.70Ta0.23O0.85C is   13     \\x0c', 'Fig. 8(d). As for the unaffected substrate, the face-centered cubic structure of the 1TaC-3HfC carbide is   almost disordered [24]. The TEM results reveal the differently contrasted areas clearly: the transition   layer is the darker one and the carbide layer is the brighter one, Fig. 8(c). The corresponding element   mapping shows how C and O vary in these two areas: the O content of the transition layer is obviously   higher than that of the substrate, whereas the C content remains stable in both areas. The HRTEM shows   the structure of the Ta-Hf-O-C oxycarbide to be clearly disordered, with neither Hf6Ta2O17 nor C   forming crystallites, and an SAED pattern that shows only diffuse rings because of short-range ordering,   composition of the substrate can be calculated as Ta0.26Hf0.74C, which agrees well with the designed   easily identified, Fig. 8(e). The HRTEM image shows that the interlayer distances are 2.65 Å and 2.30 Å   interestingly,   Journal Pre-proof   Many pores and micro-cracks were detected in the surface of the oxide scale, Fig.   1TaC-4HfC   Hf6Ta2O17 crystals mixed with HfO2 grains. In this regard, in contrast to the denser morphology of the   9(a). Furthermore, as the high-magnification SEM image shows, many sintered necks formed between the   oxide particles, Fig. 9(b), meaning that the oxide scale was sintered but still had relatively large channels   for the oxygen to pass through. The XRD and EDS results, Fig. 9(c), show that the layer comprised   corresponding to the (−111) and (200) faces, respectively. According to Vegard’s law, the chemical   1TaC-3HfC [25]. Finally and   there are some “hair-like”   tissues perpendicular   to   the   substrate in the interfacial area of the transition layer and the substrate, Fig. 8(f). This indicates that the   oxygen diffusion was clearly directional and that O diffused mainly inward along the direction of the   vertical part and reacted with the substrate Ta-Hf-C ternary carbide to form Ta-Hf-O-C oxycarbide.   pure Hf6Ta2O17 scale of 1TaC-3HfC, we infer that the difficult sintering characteristics of HfO2 hindered   the densification process of the mixed oxide layer.   14       \\x0c', 'An overview of the oxidized cross section shows the scale with a fully oxidized thickness of   approximately 250 μm, Fig. 9(d). The oxide scale looks more damaged because of the existence of the   bigger voids, indicating also that the oxide layer was quite porous, Fig. 9(e). As for the interface between   the oxide scale and the substrate, although the oxide layer was not peeled off from the substrate because   of the small RPB (RPB = 1.46), continuous intergranular cracks were found on the residual 1TaC-4HfC   grain boundary, Fig. 9(f). Furthermore, most of the fragmented grains were wrapped with a gray outer rim,   which can be considered to be similar to the oxycarbide product seen in the 1TaC-3HfC sample. This   shows that the oxide grows preferentially along the 1TaC-4HfC grain boundaries, and the mechanism of   oxide growth on 1TaC-4HfC appears to be rapid attack at the grain boundaries followed by slow   oxidation of the substrate from the grain-boundary surfaces inward.   4. Discussion   The oxidation mechanisms   in   the   Journal Pre-proof  The isothermal oxidation studies in Section 3.2 showed that the Ta-Hf-C ternary ceramics are   investigated samples can be understood   through a detailed   determine the kinetic parameters, the data in Fig. 2 for how α varies with t are analyzed further. First, note   susceptible to oxidation to varying degrees at high temperature, depending on the ir composition. To   calculation of the kinetic parameters, analysis of the reactions leading to the formation of the oxides , and   observation of the morphologies and microstructures of the oxide layers.   4.1. Oxidation kinetics   that the oxidation of Ta-Hf-C ternary ceramics is much more complicated than that of the individual   carbides. The oxidation rate of Ta-Hf-C ternary ceramics is determined by not only external factors such   as temperature but also the compositions and crystal structures of the mixed oxide layer that forms. In this   15     \\x0c', 'regard, it is not valid to characterize a complex process by using a specific parabola with only one   parabolic constant. Additionally, according to the research of Guo [26] and Inagaki [27], the oxidation   kinetics can be expressed through the exponential curve.   𝑑𝛼  𝑑𝑡  = 𝑘𝑝 (1 − 𝛼)𝑛 ,   (3)   where α is the oxidation fraction, n is the oxidation exponent, and kp is the oxidation rate constant.   Equation (3) implies that   from which kp and n can be determined from the slope and intercept of the plot of the natural logarithm of   dα/dt versus the natural logarithm of (1-α). Representative plots for the 1TaC-1HfC sample are shown in   Fig. 10(a), from which a reasonably good linear relation between ln(dα/dt) and ln(1-α) is obvious,   implying that the power law of Eq. (3) is accurate and that in principle the overall oxidation mechanism is   the same for temperatures in the range 1400-1600°C. Fig. 10(b) shows the calculated kp values for the   (4)   𝑑𝛼  𝑑𝑡  𝑙𝑛 (  ) = 𝑙𝑛 𝑘𝑝 + 𝑛 𝑙𝑛(1 − 𝛼),   Journal Pre-proof  the values of n   that.   results are 1.68 ± 0.05, 2.06 ± 0.07, 2.19 ± 0.05, and 1.87 ± 0.04   for 4TaC-1HfC, 1TaC-1HfC,   four compositions; kp increases with temperature because of the faster oxidizing reaction rate , which is   ranked as 1TaC-3HfC ＜ 1TaC-1HfC ＜ 1TaC-4HfC ＜ 4TaC-1HfC, which agrees well with the   and 1600°C obeys a rate law that is between linear and parabolic. In contrast, the values of n for   1TaC-3HfC, and 1TaC-4HfC,   respectively,   indicating   for 4TaC-1HfC and   1TaC-4HfC are both below 2.0; thus, the oxidation behavior of the two compositions between 1400°C   mass variation given in Fig. 2(a). Additionally, the value of n for each composition is obtained, and the   1TaC-3HfC are as high as 2.19, which suggests that the oxidation kinetics of 1TaC-3HfC follow a rate   law that is between parabolic and logarithmic . Conversely, the 1TaC-1HfC sample completely obeys the   16     \\x0c', 'parabolic oxidation law.   Based on   the above calculated kp results,   the apparent activation energy of Ta-Hf-C   ternary   ceramics can be calculated by the Arrhenius equation, namely   𝑘𝑝 (𝑇) = 𝐴exp (−  𝐸𝑎  𝑅𝑇  )   (5)   where, A is the pre-exponential factor, Ea is the apparent activation energy, R is the gas constant, and T is   is roughly twice those for 4TaC-1HfC and 1TaC-4HfC, making the 1TaC-3HfC good high-temperature   structural materials. By substituting   into Eq. (5),   the   temperature   dependence of kp for 4TaC-1HfC, 1TaC-1HfC, 1TaC-3HfC, and 1TaC-4HfC at 1400-1600°C can be   the temperature T for the Ta-Hf-C ternary ceramics are shown in Fig. 10(c), and the value of Ea can be   obtained from the slope of the plot. In conclusion, the obtained Ea values of the samples are summarized   in Fig. 10(d), where Ea varies an opposite tendency to kp, i.e., the highest value is that for 1TaC-3HfC,   followed by 1TaC-1HfC. In particular, the Ea value for 1TaC-3HfC is calculated to be 121 kJ/mol, which   the absolute temperature. The Arrhenius plots of the natural logarithm of kp as a function of the inverse of   Journal Pre-proof  the calculated A and Ea values   𝑘𝑝 = 12.25𝑒𝑥𝑝 (−  𝑘𝑝 = 0.60𝑒𝑥𝑝 (−  𝑘𝑝 = 0.31𝑒𝑥𝑝 (−  𝑘𝑝 = 0.42𝑒𝑥𝑝 (−  5.69 × 104  1.21 × 105  (6)   (7)   (8)   (9)   ),   ).   6.42 × 104  𝑅 × 𝑇  ),   ),   𝑅 × 𝑇  4.87 × 104  𝑅 × 𝑇  described in Eqs. (6)-(9), respectively:   𝑅 × 𝑇  4.2. Oxidation mechanisms   The calculated kinetic parameters in Section 4.1 show differing composition dependence, suggesting   that the rate-controlling mechanism changes among the samples according to the composition. Kinetics   17     \\x0c', 'controlled by oxygen diffusion is believed to dominate 1TaC-1HfC because its oxidation behavior obeys   a parabolic law. Conversely, the lower values of n of 4TaC-1HfC and 1TaC-4HfC showed that the   oxidation is partially controlled by the interfacial reactions. In contrast, the kinetic curves for 1TaC-3HfC   deviate from parabolic behavior, being closer to logarithmic-parabolic behavior and implying that except   for the oxygen diffusion control, there were some additional oxygen diffusion retardation effects during   the oxidation process of these two compositions.   An oxide map is a useful tool for predicting the oxides that formed on the sample surface. A good   consume the Ta2O5 on formation of the Hf6Ta2O17 ternary oxide, which has a relatively high melting point,   [28]. Clearly, the oxidation reactions that occur in Ta-Hf-C ternary ceramics during high-temperature   treatment in air depend on the oxidation reactions of the two individual compositions and on further   interactions among their oxide products (HfO2 and Ta2O5, respectively). If not enough HfO2 is present to   example of an oxide map is the one plotted by Kriven for the HfO2-Ta2O5 system up to 3000°C, Fig. 11   First, it is necessary to clarify the various chemical reactions that occurred in the oxidation process.   Journal Pre-proof  then the excess Ta2O5 can melt at the prevailing high temperature and then form a Ta2O5-Hf6Ta2O17   eutectic layer. The eutectic composition occurs at [HfO2]≈0.33, which is exactly the content of the HfO2   generated by the 4TaC-1HfC composition during the high-temperature oxidation process. As for the   Hf6Ta2O17 and mixed HfO2-Hf6Ta2O17, respectively. The results obtained from the oxide map agree well   with our XRD results in Section 3.3.   For the system with low HfC content (4TaC-1HfC), Ta2O5 molten liquid connected the dispersed   18   1TaC-1HfC composition characterized in detail in Fig. 6, the oxidation scale should more accurately be a   hyper-eutectic oxide. Meanwhile, for 1TaC-3HfC and 1TaC-4HfC, the oxide scales comprise pure     \\x0c', 'oxidation kinetics follows a linear-parabolic rule.   complexly entangled hyper-eutectic Ta2O5-Hf6Ta2O17 oxide   Hf6Ta2O17 crystals to be an integral whole and form a eutectic layer. The Ta2O5-Hf6Ta2O17 layer plays a   key role in reducing the inward diffusion of oxygen as the thermal barrier and provides oxidation   protection for the substrate. However, because of the small number of Hf6Ta2O17 precipitates and the high   viscosity of the Ta2O5 liquid [29], the micro-pores in the oxide scale created by the escape of the gas   products cannot be healed in time, as shown in Fig. 4(e). The existence of micro-pores and the stripping   of the oxide scale may provide tunnels for oxygen diffusion and are therefore disadvantageous to the   oxidation resistance of the 4TaC-1HfC sample. Combining the above two factors, the oxygen diffusion   retardation effects of the oxide scale of the 4TaC-1HfC composition are not considerable; thus, the   layer formed because of   the   increased   When 50 mol% of HfC was added to the 1TaC-1HfC sample, a three-dimensionally continuous and   Journal Pre-proof  the Ta2O5-Hf6Ta2O17   favorable effects   improvement of   the oxidation   regarding   the   1TaC-1HfC   composition   because   of   the   formation   of   layer. Second, because of the coexistence of the Ta2O5 liquid and the three-dimensional Hf6Ta2O17   escape of gaseous product. Third, the grain-boundary melting of the Ta2O5-Hf6Ta2O17 hyper-eutectic alloy   Hf6Ta2O17 content. There are   three   resistance   of   the   framework, cracks and pores can be infilled by Ta2O5 liquid, thereby sealing them while allowing the   hyper-eutectic oxide layer. First, with its relatively high melting point of 2450°C [30], the entangled   Hf6Ta2O17 played a “framework” role during the high-temperature oxidation process and caused a particle   enhancement effect for the viscous Ta2O5, effectively improving the mechanical performance of the oxide   can largely overcome the adverse effect of grain boundaries on oxygen diffusion, thereby improving the   oxidation resistance. Above all, we conclude that in the 1TaC-1HfC composition, the formation of a   19     \\x0c', 'stable Hf6Ta2O17-Ta2O5 hyper-eutectic scale starts limiting the oxygen penetration; thus, it decelerates the   kinetics in the form of the superimposition of a parabolic term.   As the HfC content increased further in the 1TaC-3HfC sample, Ta2O5 glass disappeared while a   double oxidation layer comprising single-phase Hf6Ta2O17 full-oxidation scale and partially oxidized   Hf-Ta-O-C scale was formed. Combining our observation results in Section 3.2, a reasonably complete   description of the oxidation mechanism and the resultant oxide scales of 1TaC-3HfC can be presented.   The first step is that the oxygen diffuses inward through the Hf6Ta2O17 layer. There are two reasons why   the oxidation resistance improves in this step, namely (i) the denser formation of the Hf6Ta2O17 inner   underneath the Hf6Ta2O17 scale becomes considerably lower than ambient atmosphere (pO2=0.2 atm), and   thus the oxygen content is insufficient to form fully oxidized Hf6Ta2O17. The second step is that a small   layer limits the oxide growth rate, and (ii) the large incorporation of Ta in the Hf6Ta2O17 facilitates a   defect structure that inhibits oxygen mobility. When this step occurs, the oxygen partial pressure (pO2)   Journal Pre-proof  a disordered-oxycarbide-crystallized-oxide   concentration   at which   thought to set in when the oxide begins to crystallize because the decrease of the amorphous phase leads   should delay further oxygen diffusion because the Hf-Ta-O-C constitute a diffusion barrier that hinders   element and the appearance of an amorphous Hf-Ta-O-C layer. The dissolution of oxygen in the lattice   amount of oxygen dissolved in the 1TaC-3HfC lattice, with the evolution of free carbon and metallic   the penetration of oxygen into the residual carbides. Finally, the oxide forms when the oxygen content   ultimately   reaches   a   transition   becomes energetically favorable. The logarithmic oxidation exhibited in the oxidation kinetic analysis is   to a decrease of the effective oxidation cross-section, thereby reducing the oxidation rate constant [31].   Compared with 1TaC-3HfC, although the HfC content of 1TaC-4HfC is only slightly increased, the   20     \\x0c', 'at the grain boundaries increased with the increase of HfC content, which we showed in our previous   oxidation mechanism differs completely from that proposed for 1TaC-3HfC. With a higher HfC content,   HfO2 and Hf6Ta2O17 mixed oxide scale formed, and because of the difficult sintering characteristics of   HfO2, it can hinder the densification process of the mixed oxide layer under the present oxidation   temperatures of 1400-1600°C. Consequently, the oxide scale of this composition looks more porous.   Moreover, by contrast to the other three compositions, many continuous intergranular cracks were found   on the residual 1TaC-4HfC grain boundaries. This is due to the fact that the amount of oxygen impurities   process with large temperature gradient. In the case of a large number of some pores and cracks in the   work [9], resulting in weaker intercrystalline bonding. Consequently, the intergranular cracks are more   prone to arise in the 1TaC-4HfC composition with the highest HfC content during the thermal shock   Journal Pre-proof  the oxidation resistance of   the   grains formed after oxidation as shown in Fig. 9(f). The gray outer oxide rim of 1TaC-4HfC grains can be   considered to be similar to the oxycarbide product seen in the 1TaC-3HfC sample, which can readily   absorb large quantities of oxygen into the lattice. And thus, a single 1TaC-4HfC grain can be regarded as   a micro-reactor, which follows the parabolic kinetics or even similar to the parabolic-logarithmic kinetics   of 1TaC-3HfC. Under the combined action of the two factors mentioned above, the oxidation kinetics of   1TaC-4HfC follows a linear-parabolic rule.   21   1TaC-4HfC sample. In this case, the oxidation kinetics of 1TaC-4HfC partially obeys the linear law. On   the other hand, the reason of parabolic kine tics is related to the oxide rim @ carbide core of 1TaC-4HfC   oxide scale, the reasons why 1TaC-4HfC sample still obeys the combining rule of linear and parabolic are   as follows. on the one hand , the existence of micro-pores and the cracks of the oxide scale may provide   tunnels for oxygen diffusion and are   therefore disadvantageous   to     \\x0c', '4.3. Recommendation for designing Ta-Hf-C ternary ceramics with improved oxidation resistance   First, we have reported that upon sintering, the HfC content that is dissolved in Ta-Hf-C ternary   ceramics changes   the densification of   the   latter:   the HfC content   is prone   to   increased oxygen   contamination that inhibits densification and meanwhile reduces the intergranular bonding of Ta-Hf-C.   This causes the oxygen to diffuse preferentially along the weak grain boundaries and then destroy the   formation of a continuous transition layer (oxygen diffusion barrier). Furthermore, during the oxidation   process, excessive HfC in Ta-Hf-C systems causes residues of refractory compound HfO2, which can   hinder the densification of the oxide scale, thereby reducing the oxidation-resistance properties. In this   the large amount of Ta2O5 present on the surface is unable to protect effectively from oxygen diffusion at   However, because of the high oxygen diffusion coefficient and the high volume expansion of Ta2O5,   Journal Pre-proof  What emerges from these considerations is that we should define the TaC/HfC ratios for preparing   in most UHTCs/SiC   to   that responsible for self-healing behavior   composites [32-35], rather than the main phases. However, the advantage of Ta2O5 is that Ta contained   liquid is more viscous than conventional SiO2 glass, thereby providing a protective phase that is more   suitable Ta-Hf-C ternary ceramics that obtain the highly dense carbides as well as limit the overall loss of   refractoriness of oxidation products and promote in situ formation of a secondary oxidation barrier.   According   to   the present   results, we   reason   that Ta-Hf-C   ternary   ceramics will have better   22   high temperatures. In this case, during the formation of the oxide layer, Ta2O5 seems to be more suitable   as   liquid filler, which   is similar   case, it appears that the lower the HfC content in Ta-Hf-C ternary ceramics (i.e., the higher the TaC)   content), the better.   resistant to volatilization [35].     \\x0c', 'comprehensive properties when the TaC/HfC ratio is between 1/1 and 1/3. Of course, the oxidation   mechanisms of Ta-Hf-C ternary ceramics are a very complex matter indeed, and we do not claim that the   present study is conclusive.   5. Conclusions   ceramics were reported and analyzed at temperatures of 1400-1600°C in air under isothermal conditions.   The following conclusions can be drawn.   1) The oxidation process of Ta-Hf-C   ternary ceramics exhibit composition dependence:   the   oxidation rate kp is ranked in the order of 1TaC-3HfC ＜ 1TaC-1HfC ＜ 1TaC-4HfC ＜ 4TaC-1HfC.   law, while 1TaC-3HfC deviates from parabolic behavior being closer to logarithmic-parabolic behavior,   thereby implying better oxidation resistance. Moreover, the Ea value of 1TaC-3HfC is calculated to be   2) Compared to 4TaC-1HfC and 1TaC-4HfC, the kinetic curves for 1TaC-1HfC obeyed parabolic   In the present work, the effects of composition on the oxidation behavior of Ta-Hf-C ternary   Journal Pre-proof  3) The new three-dimensional Hf6Ta2O17-Ta2O5 hyper-eutectic scale is found to be responsible for   4) The formation of a dense pure Hf6Ta2O17 inner layer and a continuous amorphous Hf-Ta-O-C   layer exhibits relatively   low oxygen diffusivity, and subsequently   the Hf-Ta-O-C   transition layer in 1TaC-3HfC are shown to help increase the high-temperature oxidation resistance. The   the improved oxidation performance exhibited in 1TaC-1HfC and can infill cracks and pores while   allowing the escape of gaseous product.   121 kJ/mol, which is roughly twice that of 4TaC-1HfC and 1TaC-4HfC.   pure Hf6Ta2O17   transition layer delays the further diffusion of oxygen.   5) Although the oxidation mechanisms of Ta-Hf-C ternary ceramics are a very complex matter   23     \\x0c', 'indeed and we do not claim that the present study is conclusive, we believe that Ta-Hf-C ternary   ceramics will have better comprehensive properties when the TaC/HfC molar ratio is between 1/1 and 1/3.   Acknowledgments   The authors acknowledge   the financial support of Aid Program for Science and Technology   Innovative Research Team in Higher Educational Institutions of Hunan Province , Aid program for   Innovative Research Team in National University of Defense Technology, and the support from the   National Natural Science Foundation of China under Grants 51802349.   nature or kind in any product, service and/or company that could be construed as influencing the position   presented   in, or   the review of,   the manuscript entitled “Understanding   the oxidation behavior of   Ta-Hf-C ternary ceramics at high temperature”.   24   Data availability   article.   \\uf06c   \\uf06c   \\uf06c   as references [28].   Previously reported oxide map for HfO2-Ta2O5 system was used to support this study and are   The Fig. 1 to Fig. 10 and Table 1 used to support the findings of this study are included within the   Journal Pre-proof  All data included in this study are available upon request by contact with the corresponding author.   available at [DOI:10.1111/jace.16271]. These prior studies are cited at relevant places within the text   that can inappropriately influence our work, there is no professional or other personal interest of any   We declare that we have no financial and personal relationships with other people or organizations   Conflict of interest statement       \\x0c', 'Scripta Mater. 129 (2017) 94-99.   (2016) 1539-1548.   [5] C.J. Smith, X.X. Yu, Q. Guo, C.R. Weinberger, G.B. Thompson, Phase, hardness, and deformation slip   behavior in mixed HfxTa1-xC, Acta Mater. 145 (2018) 142-153.   [6] M. 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Nasiri, Sintering   solution   formation   and   characterisation of TaC, HfC and TaC-HfC fabricated by spark plasma sintering, J. Eur. Ceram. Soc. 36   [9] J. Zhang, S. Wang, W. Li, Nano-scale 1TaC-3HfC solid solution powder synthesized using a   solvothermal method and its densification, Ceram. Int. 102 (2019) 159-163.   [10] J.M. Jiang, S. Wang, W. Li, Preparation and characterization of ultra-high temperature ternary   25   [7] S.A. Ghaffari, M.A. Faghihi-Sani, F. Golestani-Fard, H. Mandal, Spark plasma sintering of TaC-HfC       \\x0c', 'ceramics at 1073-1473 K in air, Corros. Sci. 153 (2019) 327-332.   [14] M. Desmaison-Brut, N. Alexandre, J. Desmaison, Comparison of the oxidation behavior of two   1325-1334.   [15] S. Shimada, M. Inagaki, K. Matsui, Oxidation kinetics of hafnium carbide in the temperature range   of 480° to 600 oC, J. Am. Ceram. Soc. 75 (1992) 2671-2678.   ceramics Ta4HfC5, J. Am. Ceram. Soc. 99 (2016) 3198-3201.   [11] J. Zhang, S. Wang, W. Li, Consolidation and characterization of highly dense single-phase Ta-Hf-C   solid solution ceramics, J. Am. Ceram. Soc. 102 (2019) 159-163.   [12] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y.C. Zhou, Ultra-high temperature ceramics: materials   for extreme environment applications, Wiley Press, New Jersey, 2014, pp. 13-28.   [13] B. Ye, T. Wen, D. Liu, Y. Chu, Oxidation behavior of (Hf0.2Zr0.2Ta0.2Nb0.2Ti0.2)C high-entropy   dense hot isostatically pressed tantalum carbide (TaC and Ta2C) materials, J. Eur. Ceram. Soc. 17 (1997)   Journal Pre-proof  Ionescu, Significant   improvement of   the short-term high-temperature   [17] Q. Wen, R. Riedel, E.   Corros. Sci. 145 (2018) 191-198.   oxidation resistance of dense monolithic HfC/SiC ceramic nanocomposites upon incorporation of Ta,   [18] C. Zhang, B. Boesl, A. Agarwal, Oxidation resistance of tantalum carbide-hafnium carbide solid   solutions under the extreme conditions of a plasma jet, Ceram. Int. 43 (2017) 14798-14806.   [19] K. Marnoch, High-temperature oxidation-resistant hafnium-tantalum alloys, J. Metals. 17 (1965)   1225-1231.   26   [16] C.B. Bargeron, R.C. Benson, A.N. Jette, T.E. Phillips, Oxidation of hafnium carbide   in   the   temperature range 1400 to 2060°C, J. Am. Ceram. Soc. 76 (1993) 1040-1046.     \\x0c', '[20] H. Okamoto, Hf-Ta (hafnium-tantalum), J. Phase Equilib. 17 (1996) 270-270.   [21] F.M. Spiridonov, M.N. Mulenkova, V.I. Tsirel’nikov, L.N. Komissarova, Intermediate phases in the   HfO2-Ta2O5 system, Russ. J. Inorg. Chem. 26 (1981) 922-923.   [22] Z. Cai, X. Zhao, D. Zhang, Y. Wu, J. Wen, G. Tian, Q. Cao, X. Tang, L. Xiao, S. Liu, Microstructure   and oxidation resistance of a YSZ modified silicide coating for Ta-W alloy at 1800 °C, Corros. Sci. 143   [24] Q. Xie, S.N. Wosu, The effect of TaC reinforcement on the oxidation resistance of CNTs/SiC CMCs,   J. Mater. Eng. Perform. 25 (2016) 874-883.   [23] Y. Wang, X. Xiong, G. Li, H. Liu, Z. Chen, W. Sun, X. Zhao, Ablation behavior of HfC protective   coatings for carbon/carbon composites in an oxyacetylene combustion flame , Corros. Sci. 65 (2012)   (2018) 116-128.   549-555.   [25] A. Denton, N. Ashcroft, Vegard’s law, Phys. Rev. A: At. Mol. Opt. Phys. 43 (1991) 3161 -3164.   Journal Pre-proof  [27] S. Shimada, M. Inagaki, Oxidation kinetics of hafnium carbide in the temperature range of 480 to   [28] S.J. Mccormack, K.P. Tseng, R. Weber, D. Kapush, S.V. Ushakov, A. Navrotsky, W.M. Kriven,   In-situ determination of the HfO2-Ta2O5-temperature phase diagram up to 3000 ˚C, J. Am. Ceram. Soc.   45 (2007) 1058-1065.   600℃, J. Am. Ceram. Soc. 75 (1992) 2671-2678.   [26] W. Guo, H. Xiao, Mechanisms and modeling of oxidation of carbon felt/carbon composites, Carbon.   102 (2019) 4848-4861.   [29] I.G. Talmy, J.A. Kaykoski, M.M. Opeka, High   temperature chemistry and oxidation of ZrB 2   ceramics containing SiC, Si3N4, Ta5Si3, and TaSi2, J. Am. Ceram. Soc. 91 (2008) 2250-2257.   27     \\x0c', '[30] S.J. McCormack, R.J. Weber, W.M. Kriven, In-situ investigation of Hf6Ta2O17 anisotropic thermal   expansion and topotactic, peritectic transformation, Acta Mater. 161 (2018) 127-137.   [31] L. Ogbuji, M. Singh, High   temperature oxidation behavior of reaction -formed silicon carbide   ceramics, J. Mater. Res. 12 (1995) 3232-3240.   [32] J. Hua, S. Meng, X.H. Zhang, Q. Zeng, W. Xie, Oxidation of ZrB2-SiC-graphite composites under   Mater. Sci. 44 (2009) 5673-5681.   [34] J. Li, T.J. Lenosky, C.J. Forst, S. Yip, Thermochemical and mechanical stabilities of the oxide scale   of ZrB2 + SiC and oxygen transport mechanisms, J. Am. Ceram. Soc. 91 (2008) 1475-1480.   [35] E. Emily, J. Doni, L. William, Toward oxidation-resistant ZrB2-SiC ultra-high temperature ceramics,   Metall. Mater. Trans. A, 42 (2011) 878-887.   [33] C.M. Carney, Oxidation resistance of hafnium diboride-silicon carbide from 1400 to 2000°C, J.   low oxygen partial pressures of 500 and 1500 Pa at 1800 °C, J. Am. Ceram. Soc. 99 (2016) 2474-2480.   Journal Pre-proof  28         \\x0c', 'Figure Captions   Fig. 1. (a) X-ray diffraction (XRD) pattern; (b) enlarged section of diffraction patterns for 2 θ=33-36°; (c)   typical scanning electron microscope (SEM) image with corresponding energy-dispersive spectroscopy   (EDS) analysis and elemental distribution of obtained 1TaC-1HfC sample.   Journal Pre-proof  Fig. 2. (a) Oxidation fraction (α) in TaC, HfC, and Ta-Hf-C ternary ceramics during exposure at different   temperatures for 30 min. (b-e) α versus t for Ta-Hf-C ternary ceramics: (b) 4TaC-1HfC; (c) 1TaC-1HfC;   (d) 1TaC-3HfC; (e) 1TaC-4HfC.   29             \\x0c', 'Journal Pre-proof  Fig. 3. XRD spectra of Ta-Hf-C ternary ceramics upon oxidation at 1500°C for 30 min: (a) 4TaC-1HfC;   (b) 1TaC-1HfC; (c) 1TaC-3HfC; (d) 1TaC-4HfC.   30         \\x0c', 'Fig. 4. Microstructure of 4TaC-1HfC composition after oxidization at 1500°C for 30 min: (a) overall   appearance of oxide scale surface; (b) magnified SEM image of (a); (c) corresponding EDS results for   interface between oxide scale and substrate.   points 1 and 2; (d) overall appearance of cross -section; (e) magnified SEM image of oxide scale; (f)   Journal Pre-proof  31             \\x0c', 'Fig. 5. Microstructure of 1TaC-1HfC composition after oxidization at 1500°C for 30 min: (a) overall   appearance of oxide scale surface; (b) magnified SEM image of (a) ; (c) corresponding EDS results for   points 3 and 4; (d) overall appearance of cross -section; (e) magnified SEM image of the oxide scale; (f)   interface between oxide scale and substrate.   Journal Pre-proof  HRTEM image and corresponding SAED pattern of area 1; (d) HRTEM image of area 3; (e) SAED   pattern of area 3.   Fig. 6. Transmission electron microscopy (TEM) analysis results of oxide scale: (a) STEM image and   associated EDS element mapping; (b) HRTEM image and corresponding SAED pattern of area 2; (c)   32         \\x0c', 'Fig. 7. Microstructure of 1TaC-3HfC composition after oxidization at 1500°C for 30 min: (a) overall   appearance of oxide scale surface; (b) magnified SEM image of (a); (c) corresponding EDS results for   between oxide scale and substrate.   area 5; (d) overall appearance of cross-section; (e) magnified SEM image of oxide scale; (f) interface   Journal Pre-proof  33           \\x0c', 'Fig. 8. Precise character of oxycarbide transition layer: (a) wavelength-dispersive spectroscopy (WDS)   elemental line profiles for C and O across oxycarbide transition layer; (b) Raman spectrum of transition   layer; (c) STEM image and associated EDS element mapping of transition layer; (d) STEM and SAED   results for transition layer; (e) STEM and SAED results for unaffected layer; (f) STEM result for border   area of transition layer and unaffected layer.   appearance of oxide scale surface; (b) magnified SEM image of (a); (c) corresponding EDS results for   between oxide scale and substrate.   Fig. 9. Microstructure of 1TaC-4HfC composition after oxidization at 1500°C for 30 min: (a) overall   area 6; (d) overall appearance of cross-section; (e) magnified SEM image of oxide scale; (f) interface   Journal Pre-proof  temperature; (d) calculated results for apparent activation energy Ea.   Fig. 10. Kinetic calculations for Ta-Hf-C ternary ceramics exposed to 1400°C, 1500°C, and 1600°C: (a)   natural logarithm of dα/dt versus natural logarithm of (1-α) for the representative 1TaC-1HfC sample; (b)   calculated results for oxidation rate constant kp; (c) Arrhenius plots of rate constant kp versus reciprocal   34                 \\x0c', 'Fig. 11. Oxide map for HfO2-Ta2O5 system up to 3000°C based on the work of Kriven.   Journal Pre-proof  35             \\x0c', 'Table   Table 1 Densification, hardness, elastic modulus, and thermal conductivity of Ta-Hf-C ternary ceramics.   Relative   density   Microhardness   Elastic   modulus   Thermal   conductivity   Sample   (%)   (GPa)   (GPa)   (W·m−1K−1)   1TaC-3HfC   96.5   29.92 ± 2.21   1TaC-4HfC   95.9   4TaC-1HfC   99.6   1TaC-1HfC   99.2   36.71 ± 1.21   24.46   23.93   34.65   26.24   30.27 ± 1.55   35.14 ± 1.06   559.30 ± 6.50   462.50 ± 3.10   554.70 ± 8.80   436.40 ± 13.80   Journal Pre-proof  36         \\x0c']"
},{
  "_id": 283,
  "PDF": "Volatility diagram of ZrB2-SiC-ZrC system and experimental validation.pdf",
  "Text": "['O R I G I N A L A R T I C L E  Volatility diagram of ZrB2-SiC-ZrC system and experimental validation  Ying Lu1,2  |  Ji Zou3  |  Fangfang Xu1  | Guo-Jun Zhang4  1State Key Laboratory of High  Performance Ceramics and Superfine  Microstructures, Shanghai  Institute of  Ceramics, Chinese Academy of Sciences,  Shanghai, China  2University of  the Chinese Academy of  Science, Beijing, China  3School of Metallurgy and Materials,  University of Birmingham, Birmingham,  UK  4State Key Laboratory for Modification  of Chemical Fibers and Polymer  Materials,  Institute of Functional  Materials, Donghua University, Shanghai,  China  Correspondence  Fangfang Xu, State Key Laboratory of  High Performance Ceramics and Superfine  Microstructures, Shanghai  Institute of  Ceramics, Chinese Academy of Sciences,  Shanghai, China.  Email:  ffxu@mail.sic.ac.cn  and  Ji Zou, School of Metallurgy and  Materials, University of Birmingham,  Birmingham, UK.  Email:  j.zou@bham.ac.uk  and  Guo-Jun Zhang, State Key Laboratory for  Modification of Chemical Fibers and  Polymer Materials,  Institute of Functional  Materials, Donghua University, Shanghai,  China.  Email: gjzhang@dhu.edu.cn  Funding information  National Natural Science Foundation of  China, Grant/Award Number: 51272266,  51611140121; Science and Technology  Commission of Shanghai Municipality,  Grant/Award Number: 15ZR1445200;  Shanghai  technical platform for  testing  and characterization on inorganic  materials, Grant/Award Number:  14DZ2292900  Abstract  A volatility diagram of zirconium carbide (ZrC) at 1600, 1930, and 2200°C was  calculated  in  this work. Combining  it with  the  existing  volatility  diagrams  of  ZrB2  and SiC,  the  volatility  diagram of  a  ternary ZrB2-SiC-ZrC (ZSZ)  system  was  constructed  in  order  to  interpret  the  oxidation  behavior  of ZSZ ceramics.  Applying  this  diagram,  the  formation  of ZrC-corroded  and SiC-depleted  layers  and  the  oxidation  sequence  of  each  component  in ZSZ during  oxidation  and  ablation could be well understood. Most of  the predictions from the diagrams are  consistent with  the  experimental  observations  on  the  oxidation  scale  of  dense  ZrB2-SiC-ZrC ceramics/coatings and 2200°C. The reasons for  after  oxidation  at  1600°C or  ablation  at  1930  the discrepancy are also briefly discussed.  K E Y W O R D S  microstructure, oxidation, volatility diagrams, ZrB2-SiC-ZrC  Received: 7 October 2017  |  Accepted: 5 February 2018  DOI: 10.1111/jace.15497  J Am Ceram Soc. 2018;101:3627-3635.  wileyonlinelibrary.com/journal/jace  © 2018 The American Ceramic Society  |  3627  \\x0c', '1  |  I N T R O D U C T I O N  Recent  research on hypersonic flight has inspired the inter est in developing ultra-high-temperature ceramics (UHTCs).1-4 UHTCs can withstand extremely high temperenvironments4  atures  in  chemically  aggressive  and  are  commonly defined as Group IVBand VB-based transition  metal boride, carbide, and nitride ceramics with melting temperature above 3000°C and chemical stability exceeding 2000°C.5  Thermal  shock  and  ablation  resistance  under  hypersonic flowing conditions have been widely recognized  as two key life-limiting factors for TPS (thermal protection system) applications of UHTCs.5-7  Oxidation and ablation behaviors of UHTCs have been  extensively investigated over the last 10 years. Among UHTCs, zirconium diboride (ZrB2)8-13 and hafnium diboride doped with silicon carbide (SiC) are the  (HfB2)14-17  mostly studied. Liquid B2O3 produced from the oxidation of ZrB2 forms a continuous layer on the grain surface and acts as a barrier 900°C. However, the evaporation of B2O3 at above 1400°C results in a significant loss of oxidation barrier.18,19  against  further  oxygen  diffusion  below  temperatures  SiO2, ZrB2-SiC ceramics (ZS), is considered as a critical component to improve oxidation20-24/ablation25-27 resistance of UHTCs between 1200 and 1600°C. Aforementioned evapo the  passive  oxidation  product  of  SiC in  ration of  liquids  and interactions between gases  and con densed phases during oxidation could be further understood  by volatility diagrams, which shows the thermodynamically  predicted stable  solid phase,  the concomitant gaseous  spe cies, and their vapor pressures under different oxygen par tial pressures and temperatures. For example, the ZrB2 and SiC volatility diagrams calculated by Fahrenholtz28 and Han et al27 provide a reasonable explanation for  the forma tion of  a SiC-depleted layer  in the oxidation scale of ZS  ceramics. The volatility diagram of ZrC has been calculated et al.29 CO(g)  by Maitre  and CO2(g) in Zr-C-O system, however,  are  two  important  gaseous phases  their amounts  and trends with pO2 increasing Maitre’s original work. On the other hand, only the diagrams at 1600 and 1800 K were calculated,29 the volatility  have  not  been  shown  in  diagram of ZrC at a higher  temperature is desired for a bet ter  understanding  of  the  oxidation/ablation  behavior  of  UHTCs.  Prior  results  show that  by  adding ZrC into ZrB2-SiC matrix, ternary ZSZ ceramics exhibited improved mechanical and ablation resistance.30 In addition to the SiC-depleted  layer formed on the surface of ZSZ during isothermal oxidation at 1600°C, Liu et al13 observed an extra ZrC-corroded  layer in the oxidation zone of ZSZ ceramics. These phenom ena indicate the oxidation behavior of ZSZ should be differ ent  from ZS, however,  the detailed mechanism is not clear  yet.30-32  To address these concerns,  in the present work,  the ther modynamic  calculation  and  experimental  approach  are  combined to investigate the oxidation processes of a classic  ternary ZrB2-SiC-ZrC system. Volatility diagram of ZrB2SiC-ZrC system at various temperatures were established,  the predictions  are  compared with experimental  investiga tions from this work and literature.  2  |  E X P E R IM E N T A L P R O C E D U R E S  2.1  |  Raw materials  The raw materials were self-synthesized ZrB2 powders33,34 (D50=1.05 lm, purity 98%), self-synthesized ZrC powders35 (D50=0.85 lm, purity 99%), and commercial a-SiC powders (D50= 0.45 lm, purity 98.5%, Changle Xinyuan Carborundum Co. Ltd.,  Shandong, China). The  nominal  chemical  composition (number  in vol%)  is 75ZrB2-20SiC 5ZrC (marked as ZSZ).  2.2 Powder processing, sample preparation, and sintering  |  The  starting powders were ball mixed for 24 hours using  acetone  and Si3N4 balls slurry was dried by rotary  in polyethylene  jars,  the  resultant  evaporation. After  drying,  as prepared powders mixtures were poured into a graphite die rate of 10°C/min to 1600°C in a vacuum  and heated at  a  hot-pressing  furnace. ZSZ pellet was to 1900°C under  hot  pressed  at  the  same heating rate  a pressure of 30 MPa  in flowing argon and held at  this  temperature  for 1 hour  before cooling.  2.3  |  Ablation test and characterization  The cylinders (/10 mm910 mm)  for ablation test were cut  from sintered disk by wire electrode cutting. Oxyacetylene  flame was  generated  by  oxygen  and  acetylene gas with 1930°C for  stable  flow velocity  to  ablate  the  samples  to  5 minutes. The parameters for performing the oxyacetylene  torch ablation test are listed in Table 1. During the test,  the  surface  temperature  of  the  sample was  recorded  by  an  infrared thermometer  (DIAS DSR56N) with response time  of 5 ms and measurement uncertainty of 0.5%. The corre sponding  heat  flux was  obtained  using  heat  flux  sensor  (TecFront Electronics Corp. GTT-25-5000-WF)  and  data  collector  (TecFront Electronics Corp. A2470).  After oxyacetylene  torch test,  the phase  assemblage on  the  as-ablated surface was  characterized by X-ray diffrac tion (XRD) and corresponding microstructure was observed  through  scanning  electron  microscope  (SEM,  Hitachi  S-4800)  equipped  with  energy-dispersive  spectroscopy  (EDS). Then,  the samples were mounted in resin to protect  3628  |  LU ET AL.  \\x0c', 'brittle oxide layer and cut via a diamond blade. The gradi ent  features  on  the  cross-section  of  as-ablated ZSZ and  their elemental compositions were further characterized by  SEM (JEOL  7000)  equipped with wavelength-dispersive  spectroscopy (WDS).  3  |  R E S U L T S A N D D I S C U S S I O N  3.1  |  Thermodynamic diagram  Volatility diagrams  are plots of  the vapor pressure of  the  predominant gaseous  species  as  a  function of  the  equilib rium partial pressures of O2 (pO2) and temperature. Based on the NIST-JANAF database,36 the species considered in  the Zr-B-O, Si-C-O, and Zr-C-O system in this work were  listed in Table 2. The volatility diagrams of ZrC at 1600, and 2200°C were  1930,  calculated and are  shown in Fig ure 1. The  temperature  dependence  of  pO2 the oxidation of ZrC is briefly described below. One of  and  pCO for  the  oxidation reactions of ZrC to generate CO(g)  is shown as a  demonstration for  the calculations:  ZrC(cr) þ 3 2  O2 ðgÞ ¼ ZrO2 ðcrÞ þ CO(g)  (Reaction 1)  The  relationship between Gibbs  free  energy (DG0)  and  chemical equilibrium constant  (Keq)  in reaction 1 could be  described by Equation 1  ln Keq ¼ \\x00 DG RT  ¼ ln  aCOaZrO2 ðaO2 Þ3 2 aZrC  ¼ ln  pCO=ptotal  ðpO2=ptotal Þ3  2  (1)  where R is  the  gas  constant  and  T  is  thermodynamic 2200°C.  temperature, which  is  fixed  at  1600,  1930,  or  Besides, aCO, aZrO2, aO2, and aZrC indicate the activities of CO(g), ZrO2(cr), O2(g), and ZrC(cr) in reaction 1, respectively. For the solids and liquids, their activity is equal to 1, ie, aZrC=1. For the gases, their activity is equal to the ratio of partial pressure to the total pressure, ie, aCO=pCO/ptotal. The volatility diagrams of ZrB2 and SiC were calcuapproach (Figure 2A,B),  lated  by  the  same  by  consider ing/comparing  these  diagrams  together,  the  volatility  T A B L E 1  Parameters in the flowing oxyacetylene torch ablation  test  O2 gas flux (m3/h)  1.32  C2H2 gas flux (m3/h)  0.49  Diameter of nozzle (mm)  3  Distance from sample surface to nozzle (mm)  20  Heat  flux (MW/m2)  8.5  Diameter of sample (mm)  10 1930 \\x06 30  Surface temperature (°C)  T  A  B  L  E  2  S  p  c e  i  e  s  f  o  r  t  h  e  a c  l  c  u  l  a  i t  n o  o  f  t  h  e  o v  l  a  t i l i t  y  d  i  a  g  r  a  m  s  f  o  r  Z  r  B  2  ,  a   S  i  C  ,  a  d n  Z  r  C  Z  r   B  s  p  c e  i  e  s  S  i   C  s  p  c e  i  e  s  Z  r   C  s  p  c e  i  e  s  Z  r   O  s  p  c e  i  e  s  B   O  s  p  c e  i  e  s  S  i   O  s  p  c e  i  e  s  C   O  s  p  c e  i  e  s  a g  s  e  o  u  s  C  r  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  a g  s  e  o  u  s  r c  y  s  t  a  i l l  n  e  Z  r  B   O  Z  r  B  2  r Z Z  r  O  Z  r  O  2  Z  r m   Z  r  O  2  t   Z  r  O  2  c   Z  r  O  2  B  B  2  O O B  (  B  O  )  2  B  2  O  B  O  2  B  2  O  3  B  B  2  O  3  S  i   C   O  a   S  i  C  i S S  i  O  S  i  O  2  i S S  i  O  2  C  O O C  C  O  2  C  Z  r  C   O  Z  r  C  r Z Z  r  O  Z  r  O  2  Z  r m   Z  r  O  2  t   Z  r  O  2  c   Z  r  O  2  C  O O C  C  O  2  C  LU ET AL.  |  3629  \\x0c', 'diagram of 1930°C was  a  ternary  ZrB2-SiC-ZrC Figure 2D. Such  (ZSZ)  system  at  drawn  in  construction  of  Figure 2D is  only  valid when  all  condensed  phases  (eg,  ZrB2, SiC, ZrC, do not interact  and corresponding oxides)  in ZSZ system  significantly,  ie,  all  the  solids  have  unit  activity. However,  the  reactions  and  solid  solution  forma tions  among  SiO2, concern,  ZrO2, B2O3 the volatility diagram of ZSZ was  and  are  possible.  To  address  this  calculated  separately  by  FactSage  7.0  (Figure 3)  through  an Equilib module,  in which  all  the  interactions  among  the  condensed phases  in ZSZ system were  considered. By  comparing Figures 2D and  3,  apparently,  the  critical oxy gen partial pressures  for  formation of different  regimes  for  solids  are  the  same;  the  equilibrium amounts of  the main  gaseous  substances  in  Figure 3  also  follow the  similar  trend  in  Figure 2D. Hence,  the  interactions  of  individual  solid species  in ZSZ system are  insignificant.  F I G U R E 1  The volatility diagram for zirconium carbide calculated at 1600, 1930, and 2200°C. Only the gaseous phases with \\x003 Pa are shown for than 10  equilibrium partial pressures higher  simplicity [Color  figure can be viewed at wileyonlinelibrary.com]  F I G U R E 2  Calculated volatility diagrams for A, zirconium diboride; B, a-silicon carbide; C, zirconium carbide; and D, combined ZrB2-SiCZrC at 1930°C. Only the dominant gaseous species are shown in Figure 2D for simplicity [Color  figure can be viewed at wileyonlinelibrary.com]  3630  |  LU ET AL.  \\x0c', 'Thereafter, the substances produced by the oxidation of this temperature could be predicted. At 1930°C,  ZSZ at  for  ZrB2,  two regimes could be observed from Figure 2A:  \\x006 Pa, 1. When pO2<10 and the major gaseous  the  condensed phase  is ZrB2(cr), BO(g), B(g),  phases  are  (BO)2(g), and B2O3(g); \\x006 Pa, 2. When pO2>10 the and the major gaseous  condensed phase  is ZrO2(cr), are B2O3(g), (BO)2(g),  phases  BO(g), O(g), and BO2(g).  For  SiC,  three  regimes  could  be  identified  from  Figure 2B:  \\x008 Pa, 1. When pO2<10 and the major gaseous  the condensed phase is a-SiC(cr),  phases  are  Si(g),  SiO(g),  and  CO(g); 2. When 10  \\x008<pO2<10 \\x006 Pa, exist and the major gaseous phases are SiO(g) and CO  there is no condensed phase  (g). The  oxidation  of  SiC is  active  oxidation, which  causes  the  absence  of Si-contained  condensed  phases,  and this area is called SiC-depleted layer. \\x006 Pa, 3. When pO2>10 and the major gaseous phases  the  condensed  phase  is  SiO2(l) are CO(g), CO2(g), and  O(g).  Similarly,  there  are  two  regimes  for  ZrC from Fig ure 2C:  1. When  \\x008 Pa, pO2<10 and the major gaseous phases is CO(g); \\x008 Pa, 2. When gaseous  the  condensed  phase  is ZrC(cr)  pO2>10 the major  the  condensed  phase  is ZrO2(cr) are CO(g), O(g), and  and  phases  CO2(g).  According to Figures 2D and 3, and the major oxidation 1930°C  products  in  the ZrB2-SiC-ZrC ternary with the various pO2 could be predicted:  system at  \\x008 Pa, 1. When pO2<10 the condensed phases are ZrB2(cr), a-SiC(cr), and ZrC(cr), and the major gaseous phases are Si(g), SiO(g), and CO(g);  F I G U R E 3  Calculated volatility diagram for the ZSZ ternary system at 1930°C by FactSage 7.0, only the dominant gaseous  species are shown in this picture for simplicity [Color  figure can be  viewed at wileyonlinelibrary.com]  T A B L E 3  The differences between SEM observation and thermodynamic prediction  1600°C  1930°C  2200°C  pO2 (Pa)  Predicted  condensed  phased  Observed condensed phases13  pO2 (Pa)  Predicted  condensed  phased  Observed phases  pO2 (Pa)  Predicted  condensed  phased  Observed condensed phases40  <10  \\x0011  ZrB2, SiC, ZrC  The same  <10  \\x008  ZrB2, SiC, ZrC  The same  <10  \\x008  ZrB2, SiC, ZrC  The same  10  \\x0011 10  \\x0010  ZrB2, ZrO2, C  ZrB2, SiC,  ZrO2  One layer  predicted,  two  observed  10  \\x008\\x006  10  ZrB2, ZrO2  ZrB2, SiC,  ZrO2  One layer  predicted,  two observed  10  \\x008\\x007  10  ZrB2, ZrC, Si  Si not observed  ZrB2, ZrO2  ZrB2, ZrO2, C  10  \\x007\\x006  10  ZrB2, ZrC  These two layers not  observed  10  \\x0010\\x009  10  ZrB2, ZrO2, SiO2  Two layers predicted,  only one observed:  ZrO2, SiO2  >10  \\x006  ZrO2, SiO2  ZrB2, ZrO2, SiO2  One layer  predicted,  three  observed  10  \\x006\\x003  10  ZrB2, ZrO2  >10  \\x009  ZrO2, SiO2  ZrO2, SiO2  >10  \\x003  ZrO2, SiO2  ZrO2, SiO2  One layer  predicted,  two  observed  ZrO2  ZrO2  LU ET AL.  |  3631  \\x0c', '2. When  10  \\x008<pO2<10 \\x006 Pa, ZrB2(cr) and ZrO2(cr), and are SiO(g), CO(g), O(g), and CO2(g); 3. When pO2>10 \\x006 Pa, the condensed phases are ZrO2(cr) and SiO2(l), and the major gaseous phases are B2O3(g), (BO)2(g), BO(g), O(g), BO2(g), CO(g), and CO2(g).  the  condensed  phases  are  the major  gaseous  phases  Previous reports37,38 show that the oxidation of ZrC follows ZrC?ZrCxOy?ZrO2. ZrCxOy considered as an intermediate phase in the boundary groups.39 As  process  has  been  between  ZrC and  ZrO2 thermodynamic database of ZrCxOy the NIST-JANAF table, it has not been  by  a  couple  of  the  is unavailable  in  considered  in  the  thermodynamic  calculation  in  this work. Neverthe less,  considering  the  oxidation  process,  the  pO2 equilibrium pO2 for at 1930°C.  for  ZrCxOy must be generation of ZrO2, The oxidation products  lower  than the \\x008 Pa of ZSZ at  the  ie, 10  1600  and  2200°C  were  calculated  in  the  same way as listed in Table 3, experimentally13,40  together with  those  observed  at  the  same temperature.  3.2 Exploring the volatility diagrams to explain some prior experimental results  |  The static oxidation behavior of ZrB2-SiC-ZrC ceramics at 1600°C has been investigated.13 According to Liu’s results, the oxidation scale of ZSZ ceramics at 1600°C was com posed  of  three  successive  layers. A SiO2-rich outer of ZSZ, which is followed by a SiC layer  covers  the  surface  depleted layer. There is another ZrC-corroded layer  located  between  unaffected matrix  and  the  SiC-depleted  layer.  Some  discrepancies  could  be  found  from the  predictions  and experimental observation. As shown in Table 3, C(cr) has not been observed in the ZrC-corroded layer at 1600°C in Liu’s work.13 However, according to Figure 1, C(cr) favorable over a pO2 range from 10 to 10 1600°C, its existence has also been confirmed by several groups.41,42 A possible explanation is the signal of  is  \\x0011  \\x0010 Pa  at  the light  element C might be too weak to be detected by an EDS mapping adopted in Liu’s work.13 Apart from this point, the volatility diagrams conshowed good agreement with Liu’s observation. Two partially oxidized layers, ie, ZrC-corroded layer  structed here  and SiC-depleted layer,  formed successively beneath outer  SiO2, which corresponding to a pO2 10 (Table 3). At 1600°C, SiO2 started to be gener\\x0010 Pa, ated by the passive oxidation of SiC at a pO2 of 10 which protects the ZSZ ceramics by retarding the diffusion  range  from 10  \\x0011  to  \\x0010 Pa  of oxygen into the ceramic matrix.  The  ablation  behavior  of  dense ZrB2-SiC-ZrC coating of SiC-coated C/C composites at 2200°C et al.40  on  the  surface  was  investigated  by  Zhang  An  external  ZrO2  layer,  a ZrO2-SiO2 clearly recognized. In general, volatility diagram of in Zhang’s results.40 A SiC-depleted layer \\x003 Pa was pO2 ranging both prediction confirm  layer,  and an inner SiC-depleted layer  were  ZSZ fits well  predicted  in  a  from 10  \\x007  to  10  identified,  also  experiment  and  the  existence  of  a  layer  filled  with  ZrO2 the diagram, SiO2 was formed over a 2200°C. Samples in reference40 were  and  SiO2. pO2 of ablated  Based on \\x003 Pa under high-speed  10  at  air  flow during  the  ablation  test,  and  the  very  high  temperatures  and  high  gas  velocities were  two of  the main reasons  that  caused the  absence of pro tective SiO2-based ing in an external,  layer  on  outer  oxidation  scale,  result porous,  and  columnar ZrO2 there might be an inner  layer.  It  is worth mentioning  that  layer  with  Si  remaining  between  the  matrix  and  the  SiC depleted layer at a low pO2 ranging from 10 to \\x007 Pa, however, Si has not been detected in the oxidation scale at 2200°C which might be due to its narrow range of pO2.43  \\x008  10  thermodynamic  favorable  3.3 Characterization on the ZSZ ceramics ablation at 1930°C  |  The  ZrB2-SiC-ZrC ceramics were work. Figure 4 showed the morphology,  ablated  at  1930°C in  this  element,  and  phase  assemblage  on  the  as-ablated  surface. Accord ing  to  the  SEM image  of  the  ablated  surface,  SiO2  F I G U R E 4  SEM image, EDS, and XRD for  the surface of  ablated ZSZ [Color  figure can be viewed at wileyonlinelibrary.com]  3632  |  LU ET AL.  \\x0c', 'generated  on  the  oxidation  scale was  absent  during  the  oxyacetylene  torch  ablation  test  due  to  the  very  high  temperature  and  gas  velocity.  In  addition,  the  escape  of  generated  gaseous  phases  led  to  the  creation  of  bubbles  and  holes  in  the  outer  oxidation  layer.  The  EDS  and  XRD patterns  showed that  the outer oxidation layer  con sisted of m-ZrO2. SEM image of  the  cross-section  in  a  large  area  (Fig ure 5)  showed  that  there were  five  distinctive  zones  in  oxide  layer. The  top layer only contains ZrO2, see from Figure 5, this layer is weakly bonded to the inner  as we  can  oxide structures.  WDS  elemental  mapping  (Figure 6)  of  each  zone  except  the  outer  ZrO2 clearly confirmed  layer  in  the  cross-section  of  as-ablated  ZSZ  the  existence  of ZrC corroded  layer  (V)  and SiC-depleted 70 lm,  layer  (IV) with  a  thickness  of  30  and  respectively. The  successive  F I G U R E 6  SEM images of each regimes in oxidation layer and corresponding element mapping (Si, O, B, Zr, and C) by WDS [Color  figure can be viewed at wileyonlinelibrary.com]  F I G U R E 5  SEM image and corresponding element mapping (Si,  O, B, Zr, and C) by WDS in a wide area of  the cross-section of  ablated ZSZ [Color  figure can be viewed at wileyonlinelibrary.com]  LU ET AL.  |  3633  \\x0c', 'layers  of  oxide  from the  ablated  surface  to  the  ceramic  matrix  were  ZrO2, ZrO2-SiO2, ZrO2-C, and ZrB2-SiC-ZrO2. Some differences between  ZrB2-SiO2-ZrO2,  ZrB2 the  volatility  diagram and  experimental  results  could  be  found.  For  example,  there  were  five  regimes  in the oxidation layers, while only two  regimes were predicted from the diagram. Combining the  experimental  results with the volatility diagram prediction,  a possible process for describing ceramics at 1930°C is as follows: \\x008 Pa,  the  oxidation  of  ZSZ  1. When pO2<10 is stable; 2. When  the  ternary system (ZSZ system)  10  \\x008<pO2<10 \\x006 Pa, ZrC is lowed by the active oxidation of SiC,  oxidized  first  fol the favorable con dense phases are ZrB2 and ZrO2; \\x006 Pa, passive oxidation of SiC occurred 3. When pO2>10 before the oxidation of ZrB2. SiO2(l) in the outer layer was gradually lost due to the high  temperatures  and  high velocities gas ablation.  According  to  Figure 1,  C(cr)  appeared  in  layer  IV  might be generated during cooling after  ablation test,  eg,  carbon is a favorable phase in the volatility diagram of \\x0011 and 10 \\x0010 Pa. ZrC at 1600°C under a pO2 between 10 Besides, CO(g) produced by the oxidation of ZrC at low  pO2 duct,  could retard the oxidation of SiC into the  same pro and more  likely to generate products with less oxy gen,  such as C. More  importantly,  according to Figure 6,  as  the  first  phase  to  be  oxidized,  the  oxidation  of ZrC  could  reduce  the  oxygen  levels  in  layer V,  IV,  and  III,  therefore,  the SiC and ZrB2 phases temporarily protected before all  in these  layers  could  be  the  ZrC grains  had  been consumed.  4  |  C O N C L U S I O N  In this work,  the volatility diagrams of the ZrB2-SiC-ZrC and 2200°C were drawn by  (ZSZ)  system at 1600, 1930,  both the calculation via FactSage 7.0 through the Equilib rium module  and  the  combination  of  the  diagrams  of  each  component  in  this  ternary  system. The  critical  oxy gen partial pressures  for  formation of different  regimes  in  those  two diagrams  are  the  same, which means  the  inter actions  of  individual  solid  species  in  ZSZ  system are  insignificant. A series  of  predictions  from ZrB2-SiC-ZrC by the experimental  volatility  diagrams were  confirmed  investigations  on  the  static  oxidation  or  the  ablation  of  ZrB2-SiC-ZrC ceramics via oxyacetylene of ZrC-corroded and SiC-depleted layer  torch. Formation  has  been  proved  in ZSZ system during 1930°C, which  oxidation  at  1600  and  ablation  at  temporarily  protects  the  ZrB2  grains  in  these layers.  A C K N OW L E D GM E N T  We thank the financial  supports  from the National Natural  Science  Foundation  of  China  (No.  51611140121,  51272266),  the  Science  and Technology Commission  of  Shanghai Municipality (No. 15ZR1445200),  and Shanghai  technical platform for  testing and characterization on inor ganic materials (14DZ2292900).  O R C I D  Ji Zou  http://orcid.org/0000-0002-5803-8533  R E F E R E N C E S  1.  Jackson TA, Eklund DR, Fink AJ. High speed propulsion perfor mance  advantage  of  advanced  materials.  J  Mater  Sci.  2004;39:5905-5913.  2. 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},{
  "_id": 284,
  "PDF": "word.pdf",
  "Text": "[\"Surface & Coatings Technology 357 (2019) 48-56  Contents lists available at ScienceDirect  Surface & Coatings Technology  jou rna l homepage : www .e lsev ie r .com / loca te /su r fcoa t  Ablation properties of HfB2 coatings prepared by supersonic atmospheric plasma spraying for SiC-coated carbon/carbon composites  T  Kezhi Li⁎, Guanxi Liu, Yulei Zhang⁎  State Key Laboratory of Solidification Processing, Carbon/carbon Composites Research Center, Northwestern Polytechnical University, Xi'an 710072, PR China  A R T I C L E  I N F O  A B S T R A C T  Keywords: Supersonic atmospheric plasma spraying Coating Hafnium boride Carbon/carbon composite Ablation resistance  In order to improve the ablation resistance of carbon/carbon (C/C) composites, hafnium boride (HfB2) coatings were prepared on the surface of SiC-coated C/C composites by supersonic atmospheric plasma spraying (SAPS). The anti-ablation property was investigated in an oxyacetylene torch environment under different heat flux. The phase composition, surface and cross-section microstructure and ablation behavior of the prepared coated C/C composites were characterized by using X-ray diffraction and scanning electron microscopy equipped with energy dispersion spectroscopy. The HfB2 coating exhibited excellent ablation resistance under 2400 kW/m2 after ablation for 30 s and no obvious penetrated crack and critical defect could be observed in the coating. The increase of ablation rates with the formation of rugged surface structure and pores indicated that HfB2 coating suffered more serious mechanical denudation and thermochemical erosion. The formation of protective glassy Hf-O layer with low oxygen permeability, generated due to the oxidation of HfB2 phase, inhibited the inward diffusion of oxygen and improved remarkably the ablation resistance of C/C composites.  1.  Introduction  Carbon-fiber reinforced carbon (C/C) composites had the wide applications (aerospace, engine and thermonuclear fusion), due to its many remarkable characteristics of low density, high specific strength, good thermal shock resistance and so on [1-3]. Nevertheless, the poor oxidation resistance at temperatures > 773 K limits the extensive applications of C/C composites [4]. Multi-layer ceramic coatings [5-8], resisting the diffusion and penetration of oxygen and reducing the damage of carbon fiber and the reduction of pyro-carbon matrix caused by thermochemical erosion and mechanical denudation at high temperatures, are efficient technique to protect C/C composite from oxidation. Refractory metal carbides and diborides (ZrC, TaC, HfB2) are applicable for coatings on account of their excellent performance at high temperatures [9-11]. Among the refractory metal compounds, HfB2 has received considerable attention due to its remarkable properties, such as high melting point (3523 K), high hardness and good chemical and physical stability at high temperatures. HfB2 has been considered as a promising candidate of coating material for C/C composites to maintain its mechanical properties and thermal stability at ultra-high temperature combustion environment [12-15]. Ren [13] prepared HfB2-SiC coating by in-situ reaction method and the coating could protect C/C composites from oxidation for 265 h with only 0.41 × 10−2 g/cm2  weight loss. M. Pavese [14] synthesized the HfB2/SiC coatings through painting slurry on the Cf/SiC composites. A nanocrystalline structure HfB2 coating, which exhibited a respectable hardness of 20 GPa, has ever been prepared through CVD by S. Jayaraman [16]. Nevertheless, the long fabrication period and expensive cost of in-situ reaction method and CVD technique become their own limitations. The thickness and continuity of HfB2-SiC coatings prepared by Ren [13] and Wang [17,18] through in-situ reaction method still have the room for improvement. Supersonic plasma spraying (SAPS) is a high efficient and low-cost technique for the preparation and fabrication of coatings compared with other techniques. In SAPS process, the temperature of jet arc could reach approximately 10,000 K and the velocity of the spraying particles could be up to 500 m/s [19]. The spraying particles with high speed impact on the surface of C/C composite combined with matrix compactly and arranged densely. In comparison with long preparation period of CVD and in-situ reaction method, the whole process of SAPS from spraying to the solidification of particles on the substrate and taking shape layer cost only a couple of minutes [19]. Meanwhile, the composition ratio and the thickness of coating could be controlled through adjusting spraying parameters. Yao [20] prepared ZrB2-based coating for C/C composites by SAPS and investigated the ablation resistance of the coatings. After exposed at oxyacetylene torch, the  ⁎ Corresponding authors. E-mail addresses: likezhi@nwpu.edu.cn (K. Li), 2016200548@mail.nwpu.edu.cn (G. Liu), zhangyulei@nwpu.edu.cn (Y. Zhang).  https://doi.org/10.1016/j.surfcoat.2018.09.060 Received 9 May 2018; Received in revised form 24 September 2018; Accepted 25 September 2018 Available online 26 September 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.  \\x0c\", 'K. Li et al.  Fig. 1. HfB2 particles after prilled by vacuum spray drying.  Table 1 Details of the spraying parameters for HfB2 coating.  Content  Spraying current (A) Spraying voltage (V) Primary gas Ar (L/min) Carrier gas Ar (L/min) Second gas H2 (L/min) Powder feed rate (g/min) Spraying distance (mm) Nozzle diameter (mm)  Parameters  400-430 100-150 75-100 10 5 20 100 5.5  Table 2 Ablation test parameters of the oxyacetylene torch.  Heat flux (kW/m2)  Flux (L/s)  Pressure (MPa)  Ablation angle (°)  2400 4200  C2H2  0.244 0.42  O2  0.167 0.31  C2H2  0.4 0.4  O2  0.095 0.095  90 90  Fig. 2. XRD patterns of HfB2 before and after spraying: (a) before spraying; (b) after spraying.  49  Surface & Coatings Technology 357 (2019) 48-56  coating could effectively protect C/C matrix and the ablation rates was much lower than C/C composites. Furthermore, Yang [21] investigated the ablation resistance of HfC-SiC coating prepared by SAPS. The mass and linear ablation rates were approximately 0.26 mg/s and −0.38 μm/ s after ablation under 2400 kW/m2. Zhang [22] prepared the ZrB2-SiC coating by SAPS and explored its ablation resistance under two different heat flux of 2400 kW/m2 and 4200 kW/m2. Bai [23] prepared the yttria-stabilized zirconia based thermal barrier coatings (TBCs) by SAPS. Compared to the conventional atmospheric plasma sprayed TBCs (APS-TBCs), thermal shock resistance of the SAPS-TBCs improved 90% than APS-TBCs, which was proved by the water-quenching test from 1373 K into room temperature. In the present work, a double layer coating with SiC transition layer prepared by pack cementation and HfB2 outer layer fabricated by SAPS was proposed for protecting C/C composites at high temperature environment. SiC coating has been prepared as transition layer to relieve the mismatch of CTE between C/C composites and HfB2 coating. The microstructure and phase composition of the HfB2 coating were characterized. The ablation resistance of HfB2 coatings was investigated by oxyacetylene torch system.  2. Experimental  2.1. Preparation of HfB2 coating for SiC-coated C/C composites  Cylindrical C/C composites (Φ30 mm × 10 mm) with a density of 1.70 g/cm3 were served as substrates in this work. The as-prepared substrates were polished by abrasive paper and cleaned with ethanol. The transition layer of SiC coating were prepared on the surface of substrate by pack cementation with the mixed C, Si, Al2O3 powders as raw materials. The detail of the preparation technology has been reported in Ref [24]. HfB2 coating was prepared on the SiC-coated C/C samples by supersonic plasma spraying. HfB2 particles, pelleted by vacuum spray drying, were used as the precursor powders as shown in Fig. 1. The technological parameters of spraying are listed in Table 1.  2.2. Ablation test and characterization  The ablation properties of the HfB2 coated specimens have been tested with an oxyacetylene torch system in accordance with the National Standard of Ablation Test (GJB 323A-96, China) with heat flux of 2400 kW/m2 and 4200 kW/m2. The parameters of the ablation test were exhibited in Table 2. The inner diameter of oxyacetylene gun tip was 2 mm and the distance between gun tip and specimen was 10 mm. The specimens were exposed vertically to the flame for 30 s. The flame core temperature of oxyacetylene torch could reach 3273 K. The surface temperature of coatings was recorded by an infrared thermometer connecting with the test system. The linear ablation rates (Rl) and mass ablation rates (Rm) of the specimens were severally calculated through the thickness and mass change of the specimens before and after ablation. As the following equations:  (1)  (2)  The parameters of the equations, where l0 and l1 are the thickness of specimens before and after ablation, in the same way, m0 and m1 are the mass of specimen before and after ablation and the time (t), could be measured by experiment apparatus. Surface and cross-sectional morphologies of the HfB2 coating were examined by scanning electron microscopy (SEM, JSM 6460, JEOL Ltd., Mitaka, Japan) combined with energy dispersive spectroscopy (EDS,  =Rllt()/l01=Rmmt()/m01\\x0c', \"K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  Fig. 3. Surface SEM images of HfB2 coating and cross-section backscatter micrograph of coated C/C composites: (a) low magnification; (b) high magnification of (a); (c) cross-section backscatter micrograph.  Fig. 4. Macrograph of HfB2 coated C/C composites after ablation for 30 s under different heat flux: (a) under 2400 kW/m2; (b) under 4200 kW/m2 (C, center region; T, transition region).  Oxford INCA). The phase composition and crystal structure of coating were analyzed by x-ray diffraction (XRD, X'Pert PANalytical, Almelo, Netherlands).  the Pro,  3. Results and discussion  3.1. Microstructure and morphology of HfB2 coatings  Fig. 2 shows the XRD patterns of pelleted HfB2 spray powders (before spraying) and HfB2 coating prepared by SAPS (after spraying). The  single HfB2 peaks (No. 03-065-8678) were detected in the pattern of pelleted HfB2 spray powders, which indicated that the spray powders were consisted of HfB2 particles and ensured the purity. Lower HfO2 peaks (No. 00-034-0104) were detected in the pattern of HfB2 coating proved that a little part of HfB2 particles have been oxidized during the spraying process with instantaneous high temperature environment. Conversely, B2O3 was not detected in the XRD pattern due to the rapid evaporation of B2O3 at high temperature (> 1330 K) [25,26]. The surface and cross-section SEM images were displayed in Fig. 3. As shown in Fig. 3(a) and (b), the surface, where part of spray powder  50  \\x0c\", 'Surface & Coatings Technology 357 (2019) 48-56  transition region (T). It can be seen that center region suffered the most severe ablation and a layer of white homogenous oxide-film formed on the surface. Compared the macrograph of coatings ablation under different heat flux, it could be found that the sample ablated under 4200 kW/m2 contain a larger and severer ablation center region than under 2400 kW/m2. As illustrated in Fig. 12, the initial temperature of ablation under 4200 kW/m2 is higher 200 K than under 2400 kW/m2 and the finishing temperature after ablation 30 s under 4200 kW/m2 is higher about 400 K than under 2400 kW/m2, which inferred that the coating ablation under 4200 kW/m2 suffered severer thermochemical erosion. On the other hand, the flux of C2H2 and O2 under 4200 kW/m2 is larger than under 2400 kW/m2 (Table 2), which indicated that the coating ablation under 4200 kW/m2 withstand sharper mechanical scouring and denudation. In addition, the coatings after ablation kept good integrity and no layer peeled off from the substrates, which inferred that HfB2 coating could remarkably withstand the thermochemical erosion and mechanical denudation at the high temperature combustion environment. Fig. 5 exhibited the linear and mass ablation rates of HfB2 coated C/ C composites for 30 s under different heat flux. It can be observed that HfB2 coated C/C composites presented good ablation resistance. In heat flux of 2400 kW/m2, both the linear and mass ablation rates are negative, which indicates the thickness increasing of the coating. During the ablation process, the specimens suffered thermochemical erosion, mechanical denudation. In this period, the oxidation of coating, introducing oxygen atom into the chemical compound, will lead to the thickness increasing and mass gain. Conversely, mechanical denudation and evaporation reduced the thickness and mass by remove the portion of structure of coatings. So, the thermochemical erosion was the main ablation mechanism under 2400 kW/m2 and the impact of mechanical denudation to the coating was relatively smaller than thermochemical erosion. Compared the ablation rates under 2400 kW/m2 to 4200 kW/ m2, the linear ablation rates increased from −0.32 μm/s to −0.13 μm/s and the mass ablation rates increased from −0.14 mg/s to 0.26 mg/s. These resultants demonstrate that the increase of heat flux during ablation could aggravate the thermochemical erosion and mechanical denudation. In addition, the mass ablation rate under 4200 kW/m2 indicated that the mass loss due to the mechanical denudation was more than the mass gain due to the oxidation. Thus, the effect of combustion flame scouring to the coating under 4200 kW/m2 was much stronger than scouring under 2400 kW/m2 and the mechanical denudation became the main ablation mechanism under 4200 kW/m2. Fig. 12 illustrated the temperature rising tendency of the surface of coatings during ablation. It could be observed that ablation of HfB2 outer layer under different heat flux withstood varying degrees of thermochemical erosion with distinct temperature rising tendency. The surface temperature rose instantaneously up to ~1873 K at the initial of ablation, afterwards, maintained an approximate linearly rising with the passage of time under 2400 kW/m2. The ultimate surface temperature reached ~2273 K at the end of ablation. The transient temperature of the surface at the initial of ablation was ~2073 K under 4200 kW/m2. The linearly rising stage under 4200 kW/m2 presented similar gradient with the curves of 2400 kW/m2. The ultimate surface temperature could reach ~2623 K. Compared these two curves, the higher heat flux improved the initial surface temperature and the rising trend maintained similarity because of the same thermal diffusion and conduction in the same materials of coating. In the end, the coating, ablation under 4200 kW/m2, reached higher degree of surface temperature and suffered more serious thermochemical erosion.  K. Li et al.  Fig. 5. Linear and mass ablation rates of HfB2 coated C/C composites under different heat flux.  Fig. 6. XRD patterns of the center region of HfB2 coating after ablation for 30s under different heat flux: (a) under 2400 kW/m2; (b) under 4200 kW/m2.  was not melted completely due to their very short residence times in the plasma jet, was composed of fully molten area and insufficient molten area. The stack of un-melted particles leads to the formation of diffusion channel, thus, increase the oxygen permeability. Part of holes could be filled by melted particles with the forming of compact structure closely wrapping the substrates. Cross-section backscatter micrograph of the coating is shown in Fig. 3(c). From Fig. 3(c), it can be observed that the multilayer coating exhibits double layer structure: the SiC transition layer (~80 μm) prepared via pack cementation and the HfB2 outer layer (~120 μm) deposited via SAPS. There is no obvious crack and hole could be observed at the interface between the HfB2 layer and SiC layer, which indicates a close mechanical bonding between them. The thickness of layer and close bonding is conducive to establish a main penetration and diffusion barrier for oxygen, thus, protecting the C/C substrate from oxidation.  3.2. Ablation properties  The macrograph of specimens after exposed to oxyacetylene torch vertically for 30 s under different heat flux was shown in Fig. 4. The coating surface was divided into two regions: center region (C) and  51  \\x0c', 'K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  Fig. 7. Surface morphology of the center region of HfB2 coating after ablation for 30s under different heat flux: (a) low magnification under 2400 kW/m2; (b) high magnification under 2400 kW/m2; (c) low magnification under 4200 kW/m2; (d) high magnification under 4200 kW/m2.  3.3. Ablation morphology  Fig. 6 shows the XRD pattern of HfB2 coatings after ablation for 30s in different heat flux. It could be observed that HfO2 peaks can be seen according to JCPD card No.00-034-0104. More intense and sharper diffraction peaks of HfO2 were detected compared with the XRD pattern (Fig. 2(b)) of HfB2 coating before ablation, thereby, the HfO2 phase became the main composition after ablation. The reason is that HfB2 have been oxidized dramatically and radically during the ablation test, in the meantime, one of oxidation product, B2O3, gasified due to the effect of high temperature environment [22]. From the contrast pattern of Fig. 6, it can be found that there is no obvious distinction under different heat flux, which indicates that HfB2 coatings after ablated under different heat flux generated same phase composition. Fig. 7 exhibits the microstructures of the surface of HfB2 coatings after ablation test for 30s under different heat flux. As Fig. 7 shown, some cracks could be found on the surface under the heat flux of 2400 kW/m2. The formation of cracks may be resulted from the volume contraction of HfO2 during the process of temperature reduction. Compared with the thermal expansion coefficient of SiC (4.45 × 10−6 °C−1), the coefficient of HfO2 (5.8 × 10−6 °C−1) was large and thus larger degree of volume contraction of HfO2 result in the generation of cracks [27,28]. Meanwhile, the tetragonal HfO2 would transform into monoclinic phase near 2023 K, which lead to a 2.7% expansion in unit-cell volume [29]. Rugged surface and lots of pores, as those rounds circled, were formed under the heat flux of 4200 kW/m2.  The rugged surface was caused by the combustion gas scouring of oxyacetylene torch, while, the generation and aggregation of gases caused by high-temperature oxidation leads to the formation of pores, which indicates that the mechanical denudation and thermochemical erosion under 4200 kW/m2 heat flux was much more severe than under 2400 kW/m2 heat flux. According to the previous research, materials can be sintered when the heating temperature exceed two-thirds of their melting point [30,31]. From the high magnification SEM image (Fig. 7(b) and (d)), it could be observed that a crystal layer, proved to be HfO2 by EDS and XRD, covered on the surface of the coatings, although, some of micro-holes and cracks remained due to the incompletely grown up of HfO2 grains [32]. Compared with the melting point of HfO2 (~3123 K), the surface temperature of coating could reach 2273-2623 K which is higher than two-thirds of melting point of HfO2, thus, sintering shaped HfO2 phase distributed and covered on the substrate after ablation. Due to the high melting point of HfO2, the generated crystal layer could be as a barrier inhibiting the diffusion of oxygen thereby protecting the substrate from ablation. The transition region micrographs of HfB2 coating after ablation under 2400 kW/m2 were shown in Fig. 8. In this test, the ablation temperature in center region was higher than in transition region. No obvious crack could be observed in transition region (Fig. 8(a)) compared with center region (Fig. 7(a)). From the high magnification image (Fig. 8(b)), it could be found that transition region was consists of sintered area and un-sintered area. The sintered area formed a consecutive layer covering on the surface of coating, in contrast with the  52  \\x0c', 'K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  Fig. 8. Surface morphology of the transition region of HfB2 coating after ablation for 30s under 2400 kW/m2: (a) low magnification; (b) high magnification of (a); (c) EDS analysis of Spot 1; (d) EDS analysis of Spot 2.  loosely stack of un-sintered area on the surface. It has been detected by EDS that the composition both of the two areas was Hf-O compound (Fig. 8(c) and (d)). Combined with the results of XRD patterns (Fig. 6), it could be deduced that the HfO2 did not sintered completely owing to the lower temperature under 2400 kW/m2 in transition region. The insufficient sintering and growth of HfO2 grains because of the inadequate driving force at lower temperature environment presents incompact structure in transition region, which could be the passage of oxygen diffusion and penetration thus reduce the ablation resistance of HfB2 coating. Fig. 9 displayed the micrographs of HfB2 coating of transition region after ablation under 4200 kW/m2. Compared with the transition region ablated under 2400 kW/m2 (Fig. 8), rough and fully sintered surface could be observed on the transition region ablated under 4200 kW/m2. No un-sintered Hf-O residue left on the transition region (Fig. 9(b)). The HfO2 grains grown up and sintered due to adequate driving force provided by the heat flux under 4200 kW/m2, thus an integrated Hf-O layer formed on the surface of coating. Whereas, the evaporation and aggregation of gaseous products generated at high thermal surroundings during the ablation under 4200 kW/m2 would lead to the formation of much more pinholes than under 2400 kW/m2. Fig. 10 presents the cross-section morphology and element analysis results after ablation under different heat flux. As shown in Fig. 10(a), no obvious penetrated crack could be observed in the cross-section of coating after ablation under 2400 kW/m2. However, obvious cracks  were found in the cross-section of coating after ablation under 4200 kW/m2 (Fig. 10(c)). The cracks generated due to the residual thermal stress and the volume contraction of HfO2 caused by the temperature reduction. Element analysis along with the different layer on the cross-section of Fig. 10(a) was illustrated in Fig. 10(b). Oxygen was detected in the area of HfB2 coating, which indicated the oxidation of HfB2 during the ablation. In the junction between the HfB2 coating and SiC transition layer, Si has been observed from the images, which indicated a small quantity of SiC react with the diffusing oxygen. Meanwhile, silicon diffused into HfB2 coating during ablation because selfdiffusion activation energy of silicon could be supplied at high temperature during ablation [33]. The EDS analysis to Fig. 10(c), Si has been detected in the area of HfB2 coating as shown in Fig. 10(d), demonstrated that the Si diffused into HfB2 coating through the oxidation of SiC. So Hf-O-Si layer was formed below the Hf-O layer during the ablation. It could be observed from Fig. 10(d) that HfB2 coating contains Si on the right side of the cracks while no Si could be detected on the left side of the cracks. Therefore, it could be proved that the penetrated cracks generated between the Hf-O layer and Hf-O-Si layer.  3.4. Ablation mechanism  Ablation of the HfB2 coated-C/C composites involves a variety of physical and chemical reactions, among which, mechanical denudation  53  \\x0c', 'K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  Fig. 9. Surface morphology of the transition region of HfB2 coating after ablation for 30s under 4200 kW/m2: (a) low magnification; (b) high magnification of (a); (c) EDS analysis of Area 1.  and chemical erosion are the mainly ablation mechanism under oxyacetylene torch [34,35]. In the ablation test, there are some reactions occurring on the HfB2 coating surface, as follow equations [22,36]:  (3)  (4)  (5)  (6) The ablation schematic of HfB2 coated C/C composites was illustrated as Fig. 11. Integrated and compact coating structure was presented before ablation (Fig. 11(a)). At the initial stage of ablation, the temperature of the surface of coating increased instantaneously and rapidly (Fig. 12). The impact is HfB2 would react with oxygen gradually at the oxygen-rich and high thermal temperature environment, generated HfO2 and B2O3 (Eq. (3)). Solid-state HfO2 sintered and covered on the surface of coating thereby formed a glassy layer, which is the main oxygen diffusing barrier. The evaporation and aggregation of B2O3 at high temperature resulted in the formation of pores and pinholes in the coating (Eq. (4)). Micro-cracks were propagated by the impact of CTE with the temperature increasing and the growth of oxidized layer. HfB2 coating was oxidized thoroughly due to the oxygen diffusion and infiltration inside the coating. A thin layer of Hf-O-Si barrier formed below the Hf-O layer (Eqs. (5) and (6)) because of the mild oxidation of  SiC transition layer. According to Fig. 12, the temperature degree and rising tendency under different heat flux revealed distinction. The coating ablation under 4200 kW/m2 suffered more severe mechanical sourcing inflict and more significant thermochemical erosion. Thus, the Hf-O layer would endure more fierce combustion flame impact and the ablation rates increased naturally.  4. Conclusion  HfB2 coating was prepared on the surface of SiC-coated C/C composites by SAPS. The as-prepared coating exhibits excellent ablation resistance. After exposed to combustion flame of oxyacetylene torch for 30 s under the heat flux of 2400 kW/m2, the coating protected effectively C/C matrix from ablation for 30 s. With the increase of the heat flux from 2400 kW/m2 to 4200 kW/m2, the linear and mass ablation rates of the coating increased from −0.32 μm/s to −0.13 μm/s, −0.14 mg/s to 0.26 mg/s, respectively. The center region exhibited more pinholes produced by gas evaporation and more rugged surface generated by flame scouring. The Hf-O crystal layer generated by the oxidation reaction of HfB2 during ablation, as a barrier, restrained the oxygen diffusion and reduced the oxygen permeability. The formation of cracks and pores is the main reason of the weight loss of the coated specimen ablated under 4200 kW/m2.  54  ++HfB(s)O(g)HfO(s)BO(l)22223BO(l)BO(g)2323++SiC(s)O(g)SiO(g)CO(g)22++SiC(s)O(g)SiO(g)CO(g)2\\x0c', 'K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  Fig. 10. Cross-section backscatter micrographs of coated C/C composites after ablation under different heat flux: (a) micrograph under 2400 kW/m2; (b) EDS element line analysis of (a); (c) micrograph under 4200 kW/m2; (d) spot EDS analysis of (c).  Fig. 11. Ablation schematic of HfB2 coating: (a) before ablation; (b) produce of Hf-O layer; (c) produce of Hf-O-Si  layer.  55  \\x0c', 'K. Li et al.  Surface & Coatings Technology 357 (2019) 48-56  [19]  [18]  [17]  [16]  [15]  ablation properties of carbon/carbon composites infiltrated by hafnium boride, Carbon 52 (2013) 418-426. [13] X. Ren, H. Li, Y. Chu, Q. Fu, K. Li, Ultra-high-temperature ceramic HfB2-SiC coating for oxidation protection of SiC-coated carbon/carbon composites, Int. J. Appl. Ceram. Technol. 12 (2015) 560-567. [14] M. Pavese, P. Fino, C. Badini, A. Ortona, G. 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Huang, H.J. Li, X.R. Zeng, K.Z. Li, X.B. Xiong, M. Huang, X.L. Zhang, Y.L. Liu, A new SiC/yttrium silicate/glass multi-layer oxidation protective coating for carbon/ carbon composites, Carbon 42 (2004) 2356-2359. [7] C. Isola, P. Appendino, F. Bosco, M. Ferraris, M. Salvo, Protective glass coating for carbon-carbon composites, Carbon 36 (1998) 1213-1218. F. Smeacetto, M. Salvo, M. Ferraris, Oxidation protective multilayer coatings for carbon-carbon composites, Carbon 40 (2002) 583-587. E.L. Corral, R.E. Loehman, Ultra-high-temperature ceramic coatings for oxidation protection of carbon-carbon composites, J. Am. Ceram. Soc. 91 (2008) 1495-1502. E. Wuchina, M. Opeka, S. Causey, K. Buesking, J. Spain, A. Cull, J. Routbort, F. Guitierrez-Mora, Designing for ultrahigh-temperature applications: the mechanical and thermal properties of HfB2, HfCx, HfNx, and alpha Hf(N), J. Mater. Sci. 39 (2004) 5939-5949. [11] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347-1364. [12] H. Li, D. Yao, Q. Fu, L. Liu, Y. Zhang, X. Yao, Y. Wang, H. Li, Anti-oxidation and  [10]  [8]  [9]  56  \\x0c']"
},{
  "_id": 285,
  "PDF": "Zirconia Transport by Liquid Convection during Oxidation of Zirconium Diboride-Silicon Carbide.pdf",
  "Text": "['Zirconia Transport by Liquid Convection during Oxidation of  Zirconium Diboride-Silicon Carbide  Sigrun N. Karlsdottir  w  and John W. Halloran  Department of Materials Science and Engineering, University of Michigan, Ann Arbor, Michigan 48104  Anthony Nicholas Grundy  Departement de Genie Chimique Ecole Polytechnique de Montreal, Montreal, Quebec, Canada  During high-temperature oxidation of ZrB2-SiC composites, a multi-layer oxide scale forms with a silica-rich borosilicate  liquid as the surface oxide layer. Here, a recently proposed nov el mechanism for the high-temperature oxidation of ZrB2-SiC composites is further investigated and verified. This mechanism  involves the formation of convection cells in the oxide surface  layer during high-temperature oxidation of the composite. The  formation of zirconia deposits found in the center of  the con vection cells  is proposed here to be the consequence of  liquid  transport. The nature and deposition mechanism of the zirconia  is  reported  in  detail,  using  calculated  phase  equilibrium  diagrams and microstructure observations of a ZrB2-15 vol% SiC composite tested at 15501 and 17001C in ambient air for various times. The calculated phase equilibrium diagrams for the  binary ZrO2-B2O3 system as well as the ternary B2O3-SiO2- ZrO2 system at 15001C are reported here to interpret these results.  I.  Introduction  COMPOSITES of  zirconium diboride and silicon carbide are  prominent members of the class of ultra-high-temperature ceramics (UHTC).1,2 During high-temperature oxidation of ZrB2-SiC, a complex multi-layer oxide scale forms,3,4 featuring some or all of the following features: (1) a silica-rich outer layer,  which is believed to act as a protective scale;  (2) a subscale of  crystalline zirconia, often with a columnar microstructure with  silicate between the zirconia grains; and (3) a zirconium diboride region depleted in SiC (the ‘‘SiC-depleted region’’).5 Relatively  little  boria  is  found  in  the  scale.  Boria  is  volatile  at  oxidation  temperatures,  so  it  acts  as  a  transient  liquid,  presumably evaporating from the external surface. The mecha nism of  formation of  this  complex  scale  is of  interest,  and  this paper addresses this issue from the perspective of zirconia  transport.  Recently, we  reported6  the  formation  of  distinct  flower shaped convection cells  in the  silica-rich outer  layer. These  convection cells  transported boria-rich liquid to the  surface,  where it apparently evaporated. Zirconium oxide  (ZrO2) was leading to a pattern  found in the center of the convection cells,  of ZrO2 decorating the surface. This paper reports in detail the nature and deposition mechanism of this zirconia, using  microstructure observations and calculated phase  equilibrium  diagrams.  II.  Experimental Methods  (1)  Fabrication and Oxidation Testing  The material was a hot pressed ZrB2-15-vol% SiC ceramic fabricated at ISTEC-CNR (The Institute of Science and Technol ogy for Ceramics, National Research Council, Faenza,  Italy).  Properties elsewhere.7 A ball-milled powder mixture was uniaxially hot-pressed at 18201C for 15 min at 30 MPa applied pressure.  and  processing  are  presented  in  more  detail  The tested specimens were rectangular cuboids with an average total surface area of 1 cm2. Oxidation was conducted in ambient to 15501C with a  laboratory air either by heating specimens heating rate of 131C/min and free  cooling in a conventional  furnace (Lindberg, Watertown, WI), or by the use of a novel ribbon apparatus, described in detail elsewhere.8 Briefly, the ribbon apparatus used self-supported ribbon samples,7 which  were self-heated by electrical resistance. In the conventional fur nace, the specimens were supported on pieces of the same ma terial,  and heated in stagnant  ambient  air. Specimens were  ultrasonically cleaned in acetone and dried before oxidation.  After oxidation, specimens were stored in moisture-free desic cators  to avoid any reaction of a B2O3 on the surface of specimens. Cross-section of the oxidized specimen was prepared  the  for microstructural analysis by nonaqueous polishing procedures down to 0.5-mm finish. The composition and morphology  of the oxide scales were characterized by scanning electron mi croscopy (SEM), backscattering electron microscopy (BSE), and  X-ray energy dispersive spectroscopy (XEDS) using a Philips  XL30 FEG SEM (Amsterdam,  the Netherlands). A Cameca  SX100 (Trumbull, CT) was used for electron microprobe anal ysis  (EMPA), using well-characterized mineral  standards  for  quantitative analysis of boron, zirconium, and silicon, and for  imaging in the BSE and cathodoluminscence (CL) modes.  III.  Equilibrium Phase Diagram Calculations  The phase diagrams of  the  ternary SiO2-B2O3-ZrO2 were calculated using the commercial FactSage software pack system  age. The thermodynamic properties of all solid phases are taken from the FT-Oxide database.9 The ternary liquid phase was calculated using the modified quasichemical model.10  The thermodynamic description of the liquid phase of the SiO2- B2O3 is taken from Decterov et al.11 The SiO2-ZrO2 system was optimized based on the experimental phase diagram by Butterman and Foster.12,13 These authors give the eutectic, monotectic, and critical points of the miscibility gap at 16871, 22501, and 24301C, respectively. By introducing two temperature-indepen dent positive interaction parameters, these temperatures can be reproduced within experimental accuracy giving 16881, 22521, and 23881C,  respectively. The B2O3-ZrO2 binary system was optimized based on the experimental investigations by Beard et al.14 These authors give the ZrO2 solubility as 0.9 mol% at 12001C and state that optical examination of the glass in  A. Zangvil—contributing editor  This work was supported by the Office of Naval Research (USA), under Grant No.  N 0014-02-1-0034.  w  Author to whom correspondence should be addressed. e-mail: nanna@umich.edu  Manuscript No. 23021. Received April 2, 2007; approved September 20, 2007.  Journal  J. Am. Ceram. Soc., 91 [1] 272 - 277 (2008)  DOI: 10.1111/j.1551-2916.2007.02142.x  r 2007 The American Ceramic Society  272  \\x0c', 'January 2008  Zirconia Transport during Oxidation of Zirconium Diboride-Silicon Carbide  273  Fig. 1.  Surface of a ZrB2-15 vol% SiC sample oxidized for 2 h at 15501C. (a) Back-scattered electron image showing zirconia in bright contrast; (b) same area imaged by cathodoluminesence showing petal-like lobes in dark contrast.  compositions ranging from 88 mol% B2O3-12 mol% ZrO2 to 97 mol% B2O3-3 mol% ZrO2 heated between 9001 and 14001C indicated liquid-liquid immiscibility. We optimized a single pos show either  SiO2-B2O3 Hopefully,  stable (system ZrO2-SiO2) or metastable (systems and ZrO2-B2O3) liquid-liquid miscibility gaps. experimental  investigations will  shed more  future  itive interaction parameter  for  this binary system in order to 12001C. The  solubility  at  the  reproduce  experimental ZrO2 calculated ZrO2-B2O3 phase diagram is shown in Fig. 10. The optimized phase diagram indicates a metastable miscibility gap  in the B2O3-rich region, shown as a dotted line with a critical temperature of 10621C, which could account for the liquid- liquid immiscibility observed by Beard et al.14 The liquidus of the ternary system is predicted by symmetrically extrapolating10  the binary interaction parameters to the ternary system. An iso thermal section of the ternary ZrO2-SiO2-B2O3 phase diagram, calculated at 15001C, is shown in Fig. 9. Because of the scarcity  of  be  experimental  data  the  calculated  phase  diagrams must  regarded as  somewhat  speculative. However,  the  general  features  of  the  system can  be  expected  to  be  qualitatively  correct. This ternary system is very interesting from a thermo dynamic point of view, as all  three bordering binary systems  light on this system.  IV.  Results and Discussion  Figure 1 shows the surface of the oxide film on a sample oxidized for 120 min at 15501C in a conventional furnace, sufficient to grow an oxide with about 40 mm of zirconia covered with about 20 mm of silica-rich glass. On the left  in Fig. 1(a)  is an  image taken with backscattered electrons showing the notable 100-mm  ‘‘island-in-lagoon’’  approximately  features,  where  islands in bright 500-mm diameter  contrast  are  surrounded by  approximately  ‘‘lagoons’’  in darker contrast. Note that  the  width of  these features  is  several  times  larger  than the oxide scale thickness. These features are separated by areas in lighter  gray contrast. The bright island features are shown by XEDS to  Fig. 2.  (a) A ‘‘soley’’ in the surface of ZrB2-15-vol% SiC oxidized 4 h at 15501C imaged in cathodoluminescence; (b) same soley imaged by electron probe microanalysis in (left-to right) in oxygen Ka X-rays, zirconium La X-rays, silicon Ka X-rays, boron Ka X-rays. The scale bars on the elemental maps in (b) represent the intensity of the corresponding element.  \\x0c', 'consist  of  zirconia, while  the  darker  lagoon  features  are  primarily silica. The lighter gray regions consist of silica decorated with many approximately 1-mm zirconia particles. The  same area imaged by CL is presented in Fig. 1(b) on the right.  The CL image reveals a flower-like pattern inside the lagoon  features, with petal-like lobes in dark CL contrast surrounding  the central zirconia island. These flower patterns resemble the  Islandic soley Ranuculus acris, or common Buttercup, so we will  refer to them as soleys. Figure 2 shows a soley feature in a sample oxidized at 15501C for 4 h,  imaged in CL and by X-rays  excited from oxygen (O), zirconium (Zr), silicon (Si), and boron  (B). This confirms that the lobes in dark contrast in CL are rich  in boria but relatively poor in silica, and that the central island feature is zirconia containing little silica. Elsewhere6 we have  argued that  soleys are  convection cells  that  transport a fluid  boria-silica-zirconia oxide (BSZ) liquid to the surface, where it  flows laterally, creating the lagoon features. Evaporation of the  volatile boria  from the  soley pedals  at  the  surface deposits  viscous  silica  liquid,  forming  the  lagoon features,  and solid  zirconia, forming the island features.  Zirconia is also present on the surface as discrete particles, about 1 mm in diameter. These are shown in Fig. 3, where we see  that  the small zirconia particles are not seen inside the lagoon  features, but rather between them. Notice the pattern of small  particles between the lagoon features in Fig. 3(a), which seem to  define flow patterns. The fine scale arrangement of the particles  (Fig. 3(b)), also suggests particles moved by a flowing liquid.  These fine zirconia particles can also be seen inside the glassy  silica layer, often arranged in horizontal  layers. An example is  shown in Fig. 4, a fracture surface of  the cross section of  the  glassy scale. Note the repeated layers of fine zirconia particles  separated by thin layers of particle-free silica glass. We suggest  that  these layers of zirconia particles were once decorating the  surface (as in Fig. 3), but were subsequently covered by silica rich liquid, and this process was repeated several times. The thickness of the particle-free silica layers varies from about 5 mm near the bottom of the scale to about 2 mm near the middle of  the scale, which might suggest that the extent of the lateral flow  of liquid changes with oxidation time.  In Fig. 5 we present a fractured cross section of ZrB2-15vol% SiC specimen oxidized for 15 min at 17001C in the ribbon  apparatus, showing (from the interior) unoxidized ZrB2-SiC interior, an SiC-depleted zone of ZrB2, a layer with columnar zirconia, and an outer glassy layer. The surface of the sample  shows many protuberances, which consist of a core of zirconia  surrounded by glass. The  central  zirconia ‘‘islands’’  (imaged  from the top in Figs. 1 and 2) are shown to be protuberances elevated 5-10 mm above the silicate glass  surface. We suggest  that the protuberances are associated with the convection cells.  The samples are rectangular cuboids. During oxidation the  top surface  (y-z plane)  is horizontal while  the  side  surface  (x-z plane) is vertical. Figure 6 shows a corner of a ZrB2-SiC composite oxidized for 4 h at 15501C in a conventional furnace.  The horizontal top (y-z) surface is on the left, with the vertical  (x-z) side surface to the right. Convection cells decorate the side  and top surface of the specimen. Note that the structure is sim ilar on the top and side, demonstrating that  the direction of  gravity does not influence the shape of convection features. This  suggests that the convection is not driven by density difference  (due to thermal gradient), for if it were density driven the shape  of the features would be different on the top and side, or normal vector. Elsewhere6 we  to  and  perpendicular  to  the  gravity  suggested that convection is mechanically driven by the large  molar volume increase due to oxidation. Figure 7 is an image  taken with backscattered electrons  showing a  cross-sectional  view of one of  the zirconia islands precipitated from the BSZ  Fig. 3.  Surface of specimen oxidized at 15501C for 4 hours showing features suggesting flow along the surface. (a) Shows several lagoon-like features;  note the apparent flow patterns of the fine zirconia particles between lagoons. (b) Higher magnification of the small zirconia particles between the lagoon like features highlighted by the white box in Fig. 3(a).  Fig. 4.  Fracture surface cross-section of sample showing strata of small  zirconia particles, suggesting layers of fine surface particles (as in Fig. 3)  repeatedly covered by liquid.  Fig. 5. Fractured cross-section of ZrB2-15-vol% SiC specimen oxidized for 15 min at 17001C in the ribbon apparatus showing (from the  interior): the unoxidized interior, the SiC-depleted zone, a columnar zir conia layer, and outer glassy layer.  274  Journal of the American Ceramic Society—Karlsdottir et al.  Vol. 91, No. 1  \\x0c', 'liquid when the boria was lost by evaporation. The remaining  BS liquid flowed over the surface until enough boria had evap orated to leave only remnant silica.  Details of  the growth of  the protuberances might be also  inferred from Fig. 8, which is a fracture surface of the sample oxidized at 17001C for 15 min in the ribbon apparatus. Here the  fracture plane has intersected two zirconia protuberances. The  summits of both protuberances are decorated with fine zirconia  particles. Beneath the protuberances  are what  appear  to be  columns of discrete zirconia particles in a glassy matrix. Figure  8(b)  is a higher magnification of  the columns of  the zirconia  particles, showing the structure in details. We suggest that each  of these zirconia particles was formed when the ZrO2 dissolved in the BSZ liquid precipitated at the surface. Subsequently the  zirconia particles were buried by fresh deposits from BSZ liquid,  which continued to be transported to the surface by the con vection cells.  The existence of a BSZ liquid is essential  for  the proposed  convection mechanism of  zirconia  transport.  The  system  B2O3-SiO2 is known to form liquid solutions, but experimental data for the solubility of ZrO2 in B2O3-SiO2 liquids have not been reported. The only published report of phase equilibrium in the B2O3-SiO2-ZrO2 system is by Butterman,13 but his emphasis was on the geological conditions for zircon formation.  Consequently we use computational thermochemistry to predict  the behavior of  the B2O3-SiO2-ZrO2 diagram for the isothermal  system. Figure  9  is  a  calculated  phase  section  of  the  B2O3-SiO2-ZrO2 tween a BSZ liquid solution (liquid) and crystalline phases of  system at 15001C,  showing equilibrium be zirconia (ZrO2), zircon (ZrSiO4), and silica (SiO2). We will make use of this phase diagram to interpret the phase behavior and  microstructure of the oxides in the following section. The solu bility of zirconia in a boron oxide (B2O3) system ZrO2-B2O3 is shown in Fig. 10. Notice that a transient B2O3 liquid can dissolve much more zirconia at higher temperatures. The mole fraction zirconia, XZrO2 , in the zirconia-saturated liquid at 15501C is about 0.2 in the binary, but at 18001C  liquid for the binary  the boron oxide liquid can dissolve up to 37-mol% zirconia.  This  suggests  that a transient boron oxide liquid would be a  more potent transport medium for zirconia at higher tempera tures, at least for conditions below the boiling point of boron oxide (about 22001C in air). Agents that promote boron oxide such as water,15 decrease the temperature range  vaporization,  where boron oxide liquids can exist, and should retard transport  of dissolved zirconia and silica by the transient boron oxide-rich  liquid.  (1)  Analysis of Oxide Scale Microstructure Development  Some assumptions are necessary to estimate the  composition  and amounts of  the phases produced by oxidation. It appears  that relative thickness of the SiC-depleted zone and thickness of  the zirconia layer does not change much with time. This suggests  that the diboride and the SiC oxidize in the same ratio as they  are present (no marked preferred oxidation). This implies that in  the ratio expected for oxidation of 85-vol% ZrB2-15-vol% SiC (or 79-mol% ZrB2), the composition of fresh liquid oxide should have a boron to silica ratio XB2 O3 : XSiO2 of 0.79:0.21. Solid zirconia is also produced, and some of this will dissolve in the liq uid. According to the  calculated phase diagram, a  zirconia saturated borosilicate (BSZ liquid) with this B/Si ratio dissolves  about 11 mol%.  We expect the phase assemblage for the complete oxidation of ZrB2-15-vol% SiC at 15001C to be, on a molar basis, 0.33 mole solid ZrO2 and 0.67 moles of a BSZ liquid (liquid composition 71-mol% Boria118-mol% silica111-mol% zirconia). We pre sume that the solid undissolved zirconia remains at the interface  with ZrB2 as the ‘‘primary zirconia,’’ while the BSZ liquid flows out to the surface, carrying dissolved zirconia. The BSZ liquid  can transport the dissolved zirconia to another location, where it  might precipitate as zircon or as ‘‘secondary’’ zirconia.  Boria and,  to a lesser extent,  silica can evaporate from the  BSZ liquid at the external surface, changing the composition of  the remaining liquid. Boron oxide is much more volatile, with a vapor pressure of 233 Pa at 15001C.5 The vapor pressure of silica is only 3 \\x02 10 \\x004 Pa at 15001C.16 At the surface, essentially all of the boria will evaporate. Depending on the temperature and  atmosphere, some of the silica might also evaporate, but none of  the zirconia will be removed by evaporation. Consider the case  where only the boria is volatile. As  the boria evaporates,  the  remaining liquid moves into the two-phase region liquid1solid  ZrO2, as defined by the tie lines in the phase diagram. As boria is evaporated, the remaining BSZ liquid becomes richer in silica,  and zirconia must precipitate  from the BSZ liquid. Thus we  expect the formation of secondary zirconia precipitates,  located  near the site of boria evaporation. The secondary zirconia will  be precipitated from a liquid, and might have a different mor phology from the primary zirconia, suggesting that it might be  possible to distinguish secondary zirconia within the microstruc ture. The zirconia precipitates will either remain at the location  where they formed, or be carried as a dispersed particle with the  flowing BSZ liquid.  We expect the viscosity of the BSZ liquid to change dramat ically as the fluid boria component is lost and the remaining BSZ  is enriched in silica. Very little data are available for the viscosity  of borosilicate liquids (and no data exist for zirconia-containing  borosilicate). However, we can consider the known values for the 15501C viscosity of silica (ZSiO2 5 100 GPa \\x01 s) and boria (ZB2 O3 5 40 Pa \\x01 s), combined with a viscosity for one interme Fig. 6. Corner of the surface of a ZrB2-SiC composite oxidized for 4 h at 15501C in a conventional furnace. Convection cells decorate the side  surface (x-z) and top surface (y-z) of the specimen. Note that the struc ture is similar on top and side, demonstrating that the direction of grav ity does not affect shape of the convection features.  Fig. 7. Backscattering electron microscopic image of the cross-section of a ZrB2-SiC specimen tested at 15501C for 4 h in a conventional furnace. The image shows a zirconia island along with some secondary  zirconia surfacing the oxide boria-silica-zirconia liquid.  January 2008  Zirconia Transport during Oxidation of Zirconium Diboride-Silicon Carbide  275  \\x0c', 'diate composition17 to obtain a rough estimate for the change in the borosilicate (ZBS) with mole fraction boria (XB2O3 ), which is approximately  the viscosity of  log10 ZBS ¼ 11 \\x00 9XB2O3  (1)  For a 85-vol% ZrB2-15-vol% SiC composite, tion of the fresh liquid oxide should have a boron to silica ratio  the composi XB2O3  : XSiO2 of 0.79:0.21; thus the viscosity of the boron oxiderich liquid can be roughly estimated to be around ZSiO2 \\x00B2O3 B1000 Pa \\x01 s. Eventually, as approximately 95% of the boria has liquid (ZBS) increased to the order of 10 GPa \\x01 s. evaporated, the viscosity of the remaining has If boron oxide would completely evaporate from the transient  BSZ liquid of  composition  71-mol% boria118-mol% silica  111-mol% zirconia, we would move into the two-phase region  where  zircon (ZrSiO4) precipitates phase assemblage would be a two-phase mixture,  and the final  equilibrium  zircon and  solid SiO2, in the ratio of about 75-mol% zircon-25-mol% SiO2. (We observe zirconia, not zircon, but ZrSiO4 is not very stable at these temperatures, so perhaps it does not form.) Assuming that  evaporation of boron stops due to the increase in viscosity of the  liquid when about 10-mol% B2O3 total composition moves to 58-mol% SiO2, 36-mol% ZrO2, and 6-mol% B2O3. The equilibrium phase assemblage at this composition is 66 mol% of a silica-rich liquid and 34-mol%  remains  in the  liquid,  the  solid zirconia. The microstructure of the zirconia will be different  as it is secondary zirconia precipitated from the liquid following  boron  evaporation. These  considerations  are  in  reasonable  agreement with the observed volume fractions, estimated by ste reological analysis of the glassy layer, which suggests about 75 vol% (B70 mol%) silica and about 25 vol% (B30 mol%) pre cipitated zirconia (as distinct from the columnar zirconia).  Fig. 8.  (a) Fracture surface of a ZrB2-SiC oxidized at 17001C for 15 min in the ribbon apparatus, where the fracture plane has intersected two zirconia peaks. (b) A higher magnification of the area highlighted with the white box in (a) showing the details of the zirconia peaks. Note their appearance as  columns of discrete zirconia particles in a glassy matrix.  0.1  0.2  0.3  0.4  0.5  0.6  0.7  0.8  0.9  0.1  0.2  0.3  0.4  0.5  0.6  0.7  0.8  0.9  0  .  1  0  .  2  0  .  3  0  .  4  0  .  5  0  .  6  0  .  7  0  .  8  0  .  9  ZrO2  B2O3  mole fraction  SiO2  l i  q  u  i  d  ZrSiO4  1500 °C  Fig. 9.  Calculated phase diagram for the 15001C isothermal section of  the ternary ZrO2-SiO2-B2O3, showing equilibrium between a boria-silica-zirconia liquid and crystalline phases of ZrO2, ZrSiO4, and SiO2.  Beard et al., 1962 13  Liq + c-ZrO2  Liquid+ t-ZrO2+    Liquid + m-ZrO2  Liquid  glass 1 + glass 2  mole B2O3 / (B2O3 + ZrO2)  T  (  C  )  0  .1  .2  .3  .4  .5  .6  .7  .8  .9 .9  1 1  600  800  1000  1200  1400  1600  1800  2000  2200  2400  2600  2800  Fig. 10.  Calculated binary phase diagram for the system ZrO2-B2O3, showing the solubility of zirconia in the boron oxide liquid as a function  of temperature.  276  Journal of the American Ceramic Society—Karlsdottir et al.  Vol. 91, No. 1  \\x0c', '(2)  Speculations on the Oxidation of ZrB2  It  is  tempting to examine  the microstructure of  the  zirconia  oxide layer that forms on pure zirconium diboride without ad ditives. Presumably this oxide forms crystalline zirconia and a  transient BZ liquid:  ZrB2 þ 5=2 O2 ¼ð1 \\x00 xÞðZrO2 Þcrystalline þ ðB2O3 þ x ZrO2 Þliquid  (2)  At 15001C, the binary phase diagram suggests that the mole  fraction of primary zirconia is about 0.43 of  the total compo sition, and the amount of the (B2O31ZrO2) liquid is about 0.57. The (B2O31ZrO2) liquid has a composition of 13-mol% dissolved zirconia and 87-mol% B2O3. If all the B2O3 evaporates away, the primary zirconia would be 85-mol% and the precip itated secondary zirconia 15-mol% of the total solid zirconia. Fahrenholtz16 presents a micrograph of a specimen of zirconium diboride oxidized in air at 15001C for 30 min, with about a 63-mm surface layer of porous zirconia. This particular micro graph appears to have two distinct morphologies of zirconia grains: inner grains of the order of 5-6 mm in size, perhaps with a columnar structure, and smaller grains, 1-2 mm, perhaps more  equiaxed, located closer to the external surface. It appears that the fine-grained outer zirconia layer is about 10-mm thick, while the large-grained inner layer is about 53-mm thick. This is  about the proportion one might expect if the inner grains were primary zirconia (at B85% of the total zirconia) while the outer  grains were secondary zirconia deposited after evaporation of  the boron oxide solvent. This suggests that transient liquid BZ  solution might  influence  the microstructure  of  the  porous  zirconia scales on ZrB2 even in the absence of silica from SiC.  V.  Conclusions  The formation of zirconium oxide found in the center of conintroduced in a previous article by the authors,6  vection cells,  during high-temperature oxidation of ZrB2-SiC composites, proposed here to be the consequence of liquid transport. A fluid  is  BSZ liquid is transported to the surface through the convection  cells due to a larger volume increase produced when the ZrB2- SiC material oxidizes at high temperatures, proposed in a pre vious article by the authors. There, the BSZ liquid flows later ally,  creating  convection  patterns  when  the  boria  (B2O3) center of  evaporates. Here we propose  that  the ZrO2 the cells forms when it starts precipitating from the BSZ liquid  in the  during B2O3 evaporation. When ZrO2 precipitates, becomes mainly viscous silica (SiO2) liquid, which flows outward. These conclusions are based on chemical and microstruc the liquid  tural observations as well as on phase behavior predicted by  calculated  binary ZrO2-B2O3 equilibrium diagrams. The  and  ternary B2O3-SiO2-ZrO2 calculated phase diagrams  phase  are reported here for the first time. The existence of a BSZ liq uid is essential  for the proposed convection mechanism of zir conia  transport.  The  calculated  phase  diagram for the system at 15001C  isothermal  section of  the B2O3-SiO2-ZrO2 shows equilibrium between a BSZ liquid solution (liquid) and  crystalline phases of zirconia (ZrO2), zircon (ZrSiO4), and silica (SiO2). For a ZrB2-15 vol% SiC composite (79-mol% ZrB2-21mol% SiC) we assume that the composition of the fresh liquid  oxide with a B2O3 to SiO2 ratio, XB2O3 that ratio the maximum solubility of ZrO2 in the BSZ liquid is 11 mol%, predicted from the ternary phase diagram. When the  : XSiO2 ,  is 0.79:0.21. For  B2O3 starts to evaporate from the BSZ liquid it moves into the two-phase region with solid ZrO2 and BSZ liquid. The remaining BSZ liquid then becomes richer in silica as the boria con tinuous  to evaporate, and zirconia precipitates  from the BSZ  liquid. Thus the formation of secondary zirconia located near  the site of boria evaporation is expected. From microstructure  analysis of  the surface of  the specimen the secondary zirconia  precipitates can either remain at the location where they formed  or be carried as dispersed particles by the flowing liquid, where  they serve as remnant markers of the flow pattern after the liq uid cools to the SiO2-rich glass. These findings support and expand a novel mechanism of  high-temperature oxidation of ZrB2-SiC material, based on liquid flow and formation of convection cells, proposed in a previous article by the authors.6  Acknowledgments  S. N. K. and J. W. H.  thank Carl Henderson from the EMAL for his vital  assistance with characterization, particularly CL. We thank Alida Bellosi and her  associates of ISTEC in Faenza, Italy, for providing the materials.  References  1S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem,  ‘‘Evaluation  of Ultra-High Temperature Ceramics  for Aeropropulsion Use,’’  J. Eur. Ceram. Soc., 22, 2757-67 (2002). 2M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1Hypersonic Aerosurface: Theoretical Considerations and  Historical Experience,’’ J. Mater. Sci., 39 [19] 5887-904 (2004). 3F. Monteverde and A. Bellosi,  ‘‘Oxidation of ZrB2-Based Ceramics Air,’’ J. Electrochem. Soc., 150 [11] B552-9 (2003). 4E. J. Opila, S. R. Levine, and J. Lorincz,  in Dry  ‘‘Oxidation of ZrB2and HfB2-based Ultra-High Temperature Ceramics: Effect of Ta Additions,’’ J. Mater. Sci., 39,  5969-77 (2004). 5W. G. Fahrenholtz, ‘‘Thermodynamics of ZrB2-SiC Oxidation: The Formation of a SiC-Depleted Region,’’ J. Am. Ceram. Soc., 90 [1] 143-8 (2007). 6S. N. Karlsdottir, J. W. Halloran, and C. E. Henderson, ‘‘Convection Patterns  in Liquid Oxide Films  on Zirconium Diboride-Silicon Carbide Composites  Oxidized at High Temperature,’’ J. Am. Ceram. Soc., 90 [9] 2863-7 (2007). 7S. N. Karlsdottir, J. W. Halloran, F. Monteverde, and A. Bellosi  ‘‘Oxidation  of  ZrB2-SiC: Ribbon Specimens,’’ Proceedings  Comparison  of  Furnace Heated  Coupons  and  Self-Heated  of  the  International Conference on Advanced  Ceramics and Composites, Daytona Beach, FL, January 21-26, 2007 (accepted,  May 2007). 8S. N. Karlsdottir and J. W. Halloran,  ‘‘Rapid Oxidation Characterization of  Ultra-High Temperature Ceramics,’’ J. Am. Ceram. Soc., 90 [10] 3233-8 (2007). 9C. W. Bale, P. Chartrand, S. A. Degterov, G. Eriksson, K. Hack, R. Ben Ma hfoud, J. Melanqon, A. D. Pelton, and S. Petersen,  ‘‘FactSage Thermochemical  Software and Databases,’’ Calphad, 26 [2] 189-228 (2002). Available at http://  www.factsage.com. 10A. D. Pelton and M. Blander,  ‘‘Thermodynamic Analysis of Ordered Liquid  Solutions by a Modified Quasichemical Approach—Application to Silicate Slags,’’  Metal. Trans. B, 17B, 805-15 (1986). 11S. A. Decterov, V. Swami, and I.-H. Jung, ‘‘Thermodynamic Modeling of the  B2O3, B2O3-SiO2 and the B2O3-Al2O3 Systems,’’ Int. J. Mater. Res., accepted. 12W. C. Butterman and W. R. Foster, ‘‘Zircon Stability and the Zirconium  Oxide-Silica Phase Diagram,’’ Am. Mineral., 52 [5-6] 880-5 (1967). 13W. C. Butterman,  ‘‘Equilibrium Phase Relations  among Oxides  in  the  Systems GeO2, GeO2-B2O3, HfO2-B2O3, ZrO2-SiO2-B2O3, and ZrO2-SiO2,’’ Ph.D. Dissertation, Ohio State University, Columbus, OH, 1965. 14W. C. Beard, W. C. Butterman, D. E. Koopman, H. E. Wenden, and W. R.  Foster ‘‘Research on Phase Equilibria between Boron Oxides and Refractory Ox ides, Including Silicon and Aluminum Oxides,’’ Quarterly Progress Report No. 9,  October 1, 1961-December 31, 1961, Technical Report, Ohio State University  Research Foundation, Columbus, OH, January 15, 1962. 15Q. N. Nguyen, E.  J. Opila,  and R. C. Robinson,  ‘‘Oxidation  of Ultra  High Temperature Ceramics  in Water Vapor,’’ NASA/TM-2004-212923, April  2004. 16W. G. Fahrenholtz,  ‘‘The ZrB2 Volatility Diagram,’’ J. Am. Ceram. Soc., 88  [12] 3509-12 (2005). 17R. Jabra, J. Phalippau, and J. Zarzicki, Oxide Glasses by Hot-Pressing,’’ J. Non-crystalline Solids, 42, 489-98 (1980). &  ‘‘Synthesis of Binary Glass-Forming  January 2008  Zirconia Transport during Oxidation of Zirconium Diboride-Silicon Carbide  277  \\x0c']"
},{
  "_id": 286,
  "PDF": "Zirconium carbide doped with tantalum silicide Microstructure, mechanical properties and high temperature oxidation.pdf",
  "Text": "['Materials Chemistry and Physics 143 (2013) 407e415  Contents lists available at ScienceDirect  Materials Chemistry and Physics  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t c h e m p h y s  Zirconium carbide doped with tantalum silicide: Microstructure, mechanical properties and high temperature oxidation  Laura Silvestroni a, *, Diletta Sciti a, Marianne Balat-Pichelin b, Ludovic Charpentier b  a ISTEC-CNR, Via Granarolo 64, 48018 Faenza, Italy b PROMES-CNRS, 7 rue du four solaire, 66120 Font-Romeu-Odeillo, France  h i g h l  i g h t s  \\x0f Densification of ZrC at 1970 K with addition of TaSi2. \\x0f Thorough study of the microstructure by TEM and hypothesis on the densification mechanism. \\x0f Mechanical characterization: HV e 18 GPa, KIc e 3.6 MPa m1/2, s e 500 MPa. \\x0f Oxidation tests at 1800e2200 K for 20 min and discussion on the various oxidation mechanisms.  a r t i c l e  i n f o  a b s t r a c t  A zirconium carbide ceramic was hot pressed to full density thanks to the addition of TaSi2, which enabled the densification to occur at 1970 K and improved the mechanical properties as compared to monolithic ZrC. The microstructure was analysed by combined X-ray diffraction, scanning and transmission electron microscopy to investigate the effective role of the sintering additive. In addition, high temperature oxidation was performed using the reactor REHPTS (Réacteur Hautes Pression et Température Solaire) from 1800 to 2200 K for 20 min and this composite demonstrated to resist towards the highly oxidative conditions better than other carbides, thanks to the chemical modification of the oxide formed upon Ta addition. However from 2000 K, the specimen resulted very damaged. Ó 2013 Elsevier B.V. All rights reserved.  Article history:  Received 11 January 2013 Received in revised form 12 September 2013 Accepted 18 September 2013  Keywords:  Carbides Electron microscopy Microstructure Mechanical properties Oxidation  1.  Introduction  Zirconium carbide (ZrC) is an extremely hard and refractory ceramic, belonging to the class of Ultra-High Temperature Ceramics (UHTCs) in view of its high melting point exceeding 3800 K [1]. Like all the components of this family, ZrC possesses a combination of interesting properties, like for example high corrosion resistance in both acid and basic environments, a low thermal conductivity \\x001 K \\x001) and high electrical conductivity thanks to the (20.5 W m presence of metallic bonds. Moreover, the strong covalent ZreC bonds confer to ZrC high hardness (25 GPa), high Young’s modulus (440 GPa) and mechanical strength. Being a Zr-based compound, it possesses lower density compared to other transition metal carbides, like WC, TaC or HfC. Similar to other transition metal carbides, ZrC is often sub-stoichiometric and has a stability range of  * Corresponding author. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  0254-0584/$ e see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matchemphys.2013.09.020  carbon to metal ratio from 0.65 to 0.98; when carbon content exceeds 0.98, the material contains free carbon too [2]. It is commercially used in tool bits for cutting and it is potentially suitable for re-entry vehicles, rockets or scramjet engines, where low density and high temperature load bearing capability are required. Other applications of ZrC ceramics can be found in nuclear sector, where it is used as refractory coating in reactors, thanks to its low neutron absorption cross-section and weak damage sensitivity under irradiation [3]. It has also been recognized that ZrC possesses favourable emissivity properties rendering it a promising material for use as absorber in concentrating solar power systems [4e6]. Besides these amazing chemico-physical properties, the two main drawbacks of ZrC consist in a difficult densification and a poor oxidation resistance, which characterizes all the transition metal carbides. Sintering is generally performed by pressure assisted techniques, like hot pressing (HP) [7] or spark plasma sintering (SPS) at very high temperatures [3,8,9], inducing however coarse microstructure with grain size in the order of 10e30 mm. Other  \\x0c', '408  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  pressureless sintering technologies enable the densification of ZrC at lower temperatures, but only upon addition of sintering additives such as MoSi2, ZrB2, HfC, carbon or SiC [10,11]. On the other side, poor oxidation resistance over 1070 K strongly limits the applications of ZrC and makes it necessary to work in protective environments. The final oxidation products of ZrC are ZrO2 and carbon oxides. ZrO2 forms a fine grained porous scale which allows gaseous diffusion of O2 through the pores to the ZrC surface, and therefore provides no oxidation protection. In addition, low-temperature formation of cubic modification of ZrO2, instead of thermodynamically stable monoclinic modification, may significantly contribute to the high oxidation rate as a result of about 5 orders of magnitude higher solid state oxygen diffusion in cubic zirconia compared to that in monoclinic [12]. One way to improve the oxidation resistance of ZrC is to make composites, adding for example ZrB2, SiC or silicides, able to hinder oxygen penetration through the B2O3-filled ZrO2 layer, or to form a protective silica scale [13]. In this work, TaSi2 was chosen as sintering additive for ZrC to enable densification at temperature lower than 2370 K and improve the oxidation resistance of the carbide matrix. The selection of TaSi2 was driven by its refractoriness, as it melts at 2370 K, and its proved efficacy as sintering additive for other transition metal carbides [14]. In addition, previous works on ZrB2 enabled to conclude that Ta-addition is beneficial for oxidation up to 1870 K owing to its capability to stuff oxygen vacancies in Zroxide by a higher valence cation, thus rendering the oxide layer more stable in the tetragonal crystal structure [15], or thanks to the tendency to increase the glass viscosity and induce immiscibility, which is translated in higher boiling point of the glass [16]. However tantalum becomes detrimental above 2150 K due to the melting of its oxide, Ta2O5 [17], or to the formation of a complex orthorhombic oxide, TaZr2.75O8, which has a needle-like morphology not favourable for adhesion to the unreacted bulk [15,18] and a lower melting point than pure tantalumor pure zirconium-oxides [19]. In this work, the room temperature mechanical properties of the dense ZrCeTaSi2 were measured and compared to other ZrCcompounds, then the oxidation behaviour of the ceramic was studied in the temperature range from 1800 to 2200 K in order to evaluate the performances at the different temperatures and identify eventual operating limits.  2.  Experimental procedure  2.1. Material  The ceramic was produced starting from the following commercial powders in amount 85 vol% ZrC and 15 vol% TaSi2:  cubic ZrC (H. C. Starck, Germany): Grade B, mean particle size: 3.5 mm; particle size range 0.8e6 mm; specific surface area 3 m2 g \\x001; impurities (wt%): C 1.5, O 0.6, N 0.8, Fe 0.05, Hf 2.0. hexagonal TaSi2 (Cerac Inc., Milwaukee, WI), purity 99.5%, \\x00325 mesh, particle size range 5e10 mm.  The powder mixture was ball milled for 24 h in absolute ethanol using silicon carbide milling media. Subsequently the powders were dried in a rotary evaporator, sieved through a 60-mesh screen and green shaped to 45 mm-pellet by linear pressing at 20 MPa. Hot-pressing was conducted in low vacuum (w100 Pa) using an induction-heated graphite die with a constant uniaxial pressure of 30 MPa, maximum sintering temperature of 1970 K, heating rate of \\x001 and free cooling. 20 K min  The final density was measured by Archimedes’ method on a hydrostatic balance. The theoretical density was calculated by the rule of mixtures. Crystalline phases were identified by X-ray diffraction (Siemens D500, Germany) using the CuKa radiation, with 0.04 2q-step size and 1 s scan step time. The microstructure was analysed using scanning electron microscopy (SEM, Cambridge S360) and energy dispersive spectroscopy (EDS, INCA Energy 300, Oxford Instruments, UK) on fractured and polished surfaces. TEM samples were prepared by cutting 3 mm discs from the sintered pellets. These were mechanically ground down to about 20 mm and then further ion beam thinned until small perforation were observed by optical microscope. The detailed phase analysis was performed using a transmission electron microscopy (FEI Tecnai F20 ST), with an acceleration voltage of 200 kV, equipped with an EDAX EDS X-ray spectrometer PV9761 with Super ultra-thin window. Microstructural parameters, like amount of residual porosity or secondary phases, were determined through image analysis on the micrographs of polished surfaces by commercial software (Image Pro-plus 4.0, Media Cybernetics, Silver Springs, USA).  2.2. Mechanical properties  Vickers microhardness (HV) was measured on the polished surface, with a load of 9.81 N, using a Zwick 3212 tester. Fracture toughness (KIc) was evaluated using chevron-notched 25 \\x02 2 \\x02 2.5 mm beam (CNB) in flexure. The test bars, (length \\x02 width \\x02 thickness, respectively), were notched with a 0.08 mm diamond saw; the chevron-notch tip depth and average side length were about 0.12 and 0.80 of the bar thickness, respectively. The flexural tests were performed on a semi-articulated silicon carbide four-point jig with a lower span of 20 mm and an upper span of 10 mm on a universal screw-type testing machine (Instron mod. 6025). The specimens were deformed with a cross\\x001. The slice model equation of Munz head speed of 0.05 mm min et al. was used to calculate KIc [20]. On the same machine and with the same flexural jig, the flexural strength (s), was measured on chamfered bars 25 \\x02 2.5 \\x02 2 mm (length \\x02 width \\x02 thickness, respectively), using a crosshead speed \\x001; five specimens were tested. of 0.5 mm min  2.3. High temperature oxidation resistance  Pression  et  Température  Solaire, High  (Réacteur Hautes  The reactor used to perform high temperature oxidation is the REHPTS Pressure and Temperature Solar Reactor), implemented at the focus of the Odeillo 5 kW solar furnace [21]. A heliostat reflects the incident solar flux to a concentrator with faceted mirrors. A shutter enables to control the fraction of the concentrated solar flux delivered to the sample placed inside the reactor and therefore its surface temperature. In this set-up the disc (Ø ¼ 25 mm, h ¼ 2 mm) is placed 25 mm above the focus of the solar furnace, so that very high temperatures may be achieved on a homogeneous 10 mm diameter area and at very fast rate, up to \\x001. Two mirrors enable a monochromatic (5 mm) optical 100 K s pyrometer (Ircon, Modline Plus) to measure the surface temperature of the sample through a fluorine window on a 6 mm diameter circle on the sample surface at a pyrometer-sample distance of 120 cm. To obtain the real temperature by monochromatic pyrometry, a normal spectral emissivity of 0.75 was used in these experiments, as the sample was immediately covered by an oxide layer mainly composed of zirconia. The pyrometer, together with all the parts present on the optical path, was calibrated on a blackbody.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  409  is  from  going  The accuracy of the temperature measurement 1400 \\x06 15 K to 2200 \\x06 22 K. The oxidations were performed in air with an atmosphere continuously renewed. Due to the altitude of the laboratory, the total atmospheric pressure is around 87 kPa and the oxygen partial pressure pO2 is 17 kPa. The temperature of the samples was maintained at a constant plateau during 20 min and a video camera was used to follow in situ the oxidation process. A mass spectrometer (Pfeiffer Omnistar) enabled in situ gas phase analysis. CO is expected to be one of the main gaseous products during oxidation, but its molar weight is the same as N2 one (m/e ¼ 28), so it is impossible to separate the contribution of CO from the one of preponderant N2. We therefore mainly followed the signal corresponding to m/e ¼ 44, corresponding both to CO2 and gaseous SiO. The samples were weighted before and after oxidation in order to assess the mass variation and the surfaces and cross-sections were analysed after oxidation using XRD, SEM and EDS.  3. Results and discussion  3.1. As sintered microstructure  The addition of TaSi2 to ZrC matrix enabled the achievement of \\x003, at 1970 K, which is a notably lower temfull density, 7.11 g cm perature as compared to conventional ZrC-based ceramics for which 2170e2370 K are required, see Table 1. X-ray diffraction did not evidence the presence of either ZrC or TaSi2. On the contrary a new phase with reduced cell parameter was identified as a (Zr,Ta)C solid solution (Fig. 1). The substitution of Zr atoms by Ta from TaSi2 provoked a shrinkage of the lattice from 4.692 \\x17A for ZrC to 4.646 \\x17A, which corresponds to a carbide with formula (Zr0.8Ta0.2)C, according to Vegard’s rule. The final microstructure of this composite was very complex, as the SEM images display in Fig. 2. Fig. 2a shows a homogeneous and dense microstructure with around 5 vol% of black particles of SiC. In Fig. 2b a magnification of the polished surface also shows dark grey phases with irregular faceted shape containing ZreTae SieC in different amounts, probably deriving from the original TaSi2, which is instead not clearly detected. The inset of Fig. 2b evidences the morphology of the matrix grain. As a general guideline, bright contrast indicates higher Ta content, whilst darker contrast indicates higher Zr percentage. The core with jagged edges is ZrC, moving outwards the contrast becomes brighter and EDS revealed the presence of Ta, giving rise to an approximate formula (Zr0.85Ta0.15)C. Inside the grain, 500 nm particles of ZreTaeSieC phases are visible, but the stoichiometry is difficult to define. The mean grain size of the newly formed carbide is around 5.2 mm with grains achieving up to 8 mm. Despite  the mean  being sintered at relatively low temperatures, 1970 K, grain size was notably coarsened. TEM analysis allowed a deeper investigation of the matrix: Fig. 3a evidences the coreeshell morphology with denticulate interface; the EDS spectra in Fig. 3b show the chemical composition of the two regions. Particles incorporated in SiC large grains were identified as Zr-rich silicides, Zr2Si or Zr5Si3, as those displayed in Fig. 3c. At the triple points, further ZreSi phases with flat edges were recognized, as the examples reported in Fig. 3d,e. High resolution imaging revealed the presence of dislocations between ZrC-core and (Zr,Ta)C-shell (Fig. 4a) and generally clean grain boundaries between adjacent (Zr,Ta)C grains or in contact with Zrsilicides (Fig. 4b,c) were found. Occasionally, amorphous SiO2 was noticed at the grain boundaries between the carbide and the silicide phases (Fig. 4d).  3.2. Densification behaviour  With the aim to identify the densification mechanisms activated by TaSi2, the main outcomes regarding the microstructural evolution of ZrCeTaSi2 can be summarized as follows:  ZrC original grains were surrounded by Ta-containing solid solutions, which grew epitaxially on the matrix grain (Fig. 3a). The misfit between ZrC and the (Zr,Ta)C solid solution was accommodated by corrugated grain boundaries and dislocations (Fig. 4a). These features suggest great solubility of Ta in ZrC; after sintering, no TaSi2 was observed anymore, but ZreSi phases with low dihedral angles took its place, revealing a mutual solubility of Zr in TaeSi-based phases. The new formed silicides were located mainly at the triple points (Fig. 3e). High resolution TEM showed evidence of generally clean grain boundaries between carbide matrix and silicide (Fig. 4a).  Previous studies of the interaction between carbides and TaSi2 revealed that during sintering, in presence of CO, TaSi2 tends to dissociate and decompose into Ta, gaseous SiO, liquid Si and eventually to form new TaC phase [14], according to reaction (1). The formation of liquid silicon at relatively low temperature is compatible with the enhanced sintering activity of this composite which started to shrink below 1770 K, close to the melting temperature of Si.  TaSi2 þ CO(g) / TaC þ SiO(g) þ Si(l)  (1)  Free Si was not clearly observed in the final microstructure; however in the highly reducing environment of the hot pressing chamber with graphite dies and rams, it is very probable that liquid silicon was reduced to SiC, as testified by Fig. 2:  Table 1 Sintering parameters, microstructural features and mechanical properties of some ZrC-based ceramics. Legend: HP-hot pressing, SPSspark plasma sintering, PLSepressureless sintering, m.g.s. e mean grain size, HV e hardness, KIceCNB fracture toughness, s e 4-pt flexural strength.  Label  Sintering  ZCT ZCM ZCM ZCM ZCZg ZCS  K, min, atm, MPa  HP 1970, 6, vac, 30 HP 2170, 12, vac, 30 SPS 1970, 3, vac, 100 PLS 2220, 60, Ar, HP 2170, 60, vac, 30 PLS 2370, 120, Ar,  a Direct crack measurements. b 3-pt bending.  Sint. add.  vol%  15 TaSi2 15 MoSi2 9 MoSi2 20 MoSi2 8 wt Zr þ 2.3 wt graphite 20 SiC  Rel. r  %  99.9 96.8 99.0 96.8 98.3 96.7  m.g.s  mm  5.2 3.9 3.5 6.0 w10 3.1  HV  GPa 17.9 \\x06 0.7 19.2 \\x06 0.4 20.0 \\x06 0.5 12.7 \\x06 1.0 16.2 \\x06 0.9 11.8 \\x06 0.8  KIc  (MPa m1/2) 3.6 \\x06 0.4 3.4 \\x06 0.5 3.3 \\x06 0.4a 3.5 \\x06 0.2 4.7 \\x06 0.4a  e  s  MPa 503 \\x06 55 474 \\x06 41 591 \\x06 48b 272 \\x06 12  e  474b  Ref  This work [23] [8] [11] [7] [25]  \\x0c', '410  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 1. X-ray diffraction pattern of the as sintered ZrCeTaSi2 composite showing peaks shift from the pure ZrC phase.  2Si(l) þ CO(g) / SiC þ SiO(g)  (2)  Those findings are in good accordance with the microstructural features, as pure TaSi2 is not found anymore and TaC is instead present in the solid solution with ZrC according to Eq. (3):  The well-defined boundary between core and shell and the morphology of the interface between them with trapped particles also put forward a re-precipitation from liquid phase over a diffusion process.  TaC þ ZrC / (Zr,Ta)C  3.3. Mechanical properties  (3)  Indeed, it is well known that the solubility between carbides of Group IV and mono-carbides of Group V is complete and they are expected to form solid solutions [1]. Microstructural features and thermodynamics suggest that densification occurred through transient liquid phase. This hypothesis is further strengthened by the irregular shape displayed by the residual silicide phases, as well as the relevant coarsening of the carbide grains. We can reasonably suppose that liquid phase containing TaeSieCeO, where Zr was soluble in, formed during densification and crystallized at the triple points in form of ZreSi phases upon cooling leaving clean grain boundaries. As solid solutions formed, Ta from TaSi2 substituted Zr atoms in ZrC lattice. This may occur either by cations diffusion or by solution-reprecipitation. Given the low self-diffusion coefficient of this class of materials, it is presumed that lattice diffusion can occur only at very high temperature. Indeed, solution re-precipitation seems to be the dominant mechanism, in light of the sintering behaviour characterized by a relatively low shrinkage temperature.  Table 1 summarizes the main sintering parameters, microstructural features and mechanical properties of the ZrCeTaSi2 ceramic, compared to other ZrC-based materials taken from literature. The hardness of the ZrCeTaSi2 composite was about 18 GPa, in the range of the data reported in the literature for similar materials. The discrepancy between monolithic ZrC, 25 GPa, and this composite is due to the presence of softer silicide phases which have hardness below 16 GPa [22] and the quite coarse microstructure. The fracture toughness resulted in the range or higher than the other ZrC-composites, indicating that this property is mainly dictated by ZrC matrix toughness [7,8,11,23,24]. ZrCeTaSi2 composite displayed strength around 500 MPa, higher than similar composites containing MoSi2 processed by hot pressing and pressureless sintering [7,8,11,23,24]. The highest value shown in Table 1 for the SPS-composite refers to measurements carried out on 3-point bending and very small specimens [8]. So it is concluded that this ceramic displayed very good mechanical properties also after sintering at such low temperature (1970 K).  Fig. 2. SEM images of the polished surface of ZrCeTaSi2 composite showing (a) the formation of SiC dark particles and (b) the morphology of the matrix.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  411  Fig. 3. TEM images of the as sintered ZrCeTaSi2 composite showing (a) the morphology of the matrix grain with the corresponding EDS spectra in (b), (c) SiC grains incorporating ZreSi phases, (d) and (e) ZreSi phases at the triple point.  3.4. Microstructure of oxidized specimens  Pictures of the samples after testing at 1800, 2000 and 2200 K for 20 min, are shown in Fig. 5. The front faces are covered with a grey layer becoming progressively white indicating that the expected oxidation occurred. The oxide layer that formed at 1800 K, Fig. 5a, looks quite smooth and well adherent to the carbide, indicating that at this stage the sample well survived the test. At  2000 K, Fig. 5b, the sample broke upon cooling and some bubbles can be seen in the more heated zone. At 2200 K, Fig. 5c, bubbling phenomena are more evident, the surface is whiter and a wavy surface is also noticeable, indicating a change in the oxidation resistance at this temperature. Fig. 6 shows the evolution of CO2 and SiO as a function of the oxidation temperature, determined using mass spectrometry. We can observe that at 1800 K the material evolves very low amount of  Fig. 4. HR-TEM images of the as sintered ZrCeTaSi2 composite showing (a) the interface between ZrC-core and (Zr,Ta)C shell with dislocations indicated by arrows, (b), (c) examples of clean grain boundaries between (Zr,Ta)C e (Zr,Ta)C e ZrSi2, or partially wetted grain boundaries in (d).  \\x0c', '412  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 5. Photographic images of the ZrCeTaSi2 discs after oxidation tests at (a) 1800, (b) 2000 and (c) 2200 K.  Fig. 9 presents the cross-section of the sample oxidized at 1800 K, where a 140 mm thick layer underwent oxidation with formation of the mixed Zr, Ta oxide with granular shape covered by 3 mm thin and discontinuous silica layer (Fig. 9b). Moving further inward, down to around 400 mm from the surface, a complex mixture of (Zr,Ta)Si2, (Zr,Ta)-oxy-carbide and SiC phases were found, Fig. 9c. The complex (Zr,Ta)eCeO phase is the result of partial oxidation of the starting matrix grain, made also of (Zr,Ta)C, as outlined above. The interfaces between the oxide, oxy-carbide and the bulk are crack free and continuous. The formation of an oxy-carbide standing between the pure oxide external layer and the carbide core, had been already reported for ZrC and HfCcompounds [25,26]. After oxidation at 2000 K, the specimen resulted in a layered structure with the outermost scale composed of 5 mm long TaZr2.75O8 grains standing out on a continuous silica-based scale (Fig. 10b), which topped about 300 mm of coarse ZrO2 and TaZr2.75O8 grains, partially filled with silica and where vigorous bubbling clearly occurred Fig. 10a. This thick external layer was partially detached from the underneath surface, which is not shown, since it had the same aspect as the section oxidized at 1800 K (Fig. 9). The cross-section of the ceramic oxidized at 2200 K is shown in Fig. 11, where all the thickness of the disc resulted partially oxidized. The outermost thick scale is composed by a compact ZrO2 layer where 20e30 mm large porosities can be found (Fig. 11a). In this region, magnified in Fig. 11b, SiO2 droplets (dark) containing Zr traces are surrounded by a brighter phase, identified as a solid  gaseous species, which notably increases after 10 min of oxidation at 2000 K. This trade off can be due to the migration of silicide phases to the surface, which form a protective glassy scale until the dissociation of ZreSi silicides prevails, resulting in further development of SiO. At 2200 K the gases escaping is notably higher, owing to the complete dissociation of the silica-glass and the lack of a protective oxide scale, or to CO2 escape through fissures. As far as concerns mass variation, as expected, the weight gain increased with increasing temperature going from \\x002 min \\x001 at 1800 K, to 2.37 mg cm \\x002 min \\x001 at 2000 K 1.20 mg cm \\x002 min \\x001 at 2200 K. and 3.85 mg cm X-ray diffraction patterns collected on the surface of the three specimens after oxidation tests are shown in Fig. 7. The main crystalline phases at all temperatures is a mixed oxide with composition TaZr2.75O8, which has an orthorhombic structure and preferred orientation along the [020] planes; the peaks become sharper increasing the temperature, indicating improved crystallization. Monoclinic ZrO2 is also present and the peaks signal increases with temperature to about 25%1 at 2200 K, indicating the higher stability of pure oxide over the mixed one. Cubic and tetragonal ZrO2 could be present too, but the superimposition of the main peaks with the mixed oxide hinders a conclusive analysis. Other authors [25,26] and similar tests on ZrC-based composites showed indeed that the presence of carbon, coming from the oxidation of the carbide, stabilizes c-ZrO2 at low temperatures. Fig. 8 presents SEM images of the surfaces of the central region of the discs at 1800, 2000 and 2200 K. The addition of TaSi2 to a ZrC matrix generated a variety of elaborated morphologies varying the oxidation temperature; as a rule of thumb, dark regions correspond to silica-based glass, bright phases to oxides. At 1800 K the surface of the composite presents a rough cracked aspect with TaZr2.75O8 and ZrO2 being the main phases where discrete pockets of silica-based glass are found, Fig. 8a. The surface at 2000 K is mainly composed of petal-like grains of TaZr2.75O8 which form volcanos and tend to microcracking, Fig. 8b. At 2200 K melting and recrystallization of the oxide occurred and 20 mm large grains formed leaving residual silica, containing Zr traces, at the grain boundaries, Fig. 8c. These grains are in turn composed of polyhedral structures of ZrO2 and TaZr2.75O8 showing the growth planes decorated by a dendritic irregularly shaped phase identified as ZrO2 and ZrSiO4 containing small traces of Ta, inset in Fig. 8c.  1  Since the scattering coefficient for the orthorombic TaZr2.75O8 phase has not been published in the literature and is not available in the ICSD database, we considered this phase as a solid solution between 0.5 mol orthorombic Ta2O5 (#54e 514) and 3 mol tetragonal ZrO2 (#42.1164) and estimated a scattering coefficient of 4.7.  Fig. 6. Measured concentration of SiO and CO2 (m/e ¼ 44) produced at (a) 1800, (b) 2000 and (c) 2200 K during the oxidation of ZCT.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  413  Fig. 7. X-ray diffraction patterns of the ZrCeTaSi2 composite after oxidation at (a) 1800, (b) 2000 and (c) 2200 K.  solution with possible formula (Zr0.8Ta0.2)O2. Right underneath this 550 mm layer, the mixed TaZr2.75O8 oxide is present in form of a thick dense scale including elongated porosity and a Ta-rich white oxide (Fig. 11c, upper part). Moving further inward, fine grained ZrO2 with Ta traces and SiO2 discrete phases (Fig. 11d) stand above the already mentioned (Zr,Ta)-oxy-carbide and SiC phases (Fig. 11c lower part and Fig. 11e). The core of the disc is composed by a mixed (Zr,Ta)-oxy-carbide containing progressively lower oxygen amounts. These articulate morphologies are the result of complex oxidation mechanisms, including melting, phase separation and reprecipitation, occurring at 2200 K.  3.5. Oxidation behaviour  Pictures and SEM images clearly reveal this composite to undergo various oxidation mechanisms depending on the different temperature range. Fig. 6 evidences that the evolution of CO2 and SiO gaseous species has a slow and parabolic trend at 1800 K for the whole duration of the test, but at 2000 K an abrupt increase occurs after 10 min of oxidation. This variation was ascribed to the melting and decomposition of Zr-silicides, occurring at around 1900 K [27], which resulted in further SiO release. A second explanation could  be the migration of silica glass to the surface, as observable in Fig. 10b, which after a certain period loses its shielding action and is no longer protective to gases escape, as indicated by the large voids in the cross-section of Fig. 10a, and by the bubbling and fracture of the sample in Fig. 5b. From this temperature on, vigorous gas escape was reported to occur, as demonstrated also by the turbulent microstructure in Fig. 11a. It is almost established that the oxidation of carbides of groups IVeVI transition metals occurs through formation of an oxy-carbide of the metal plus carbon, which is subsequently oxidized to CO and CO2, ending with formation of the metal oxide [25,26], according to reactions of the type:  MeC þ 2O2 / MeCxOy þ C1\\x00x þ O4\\x00y / MeO2 þ CO2  (4)  This general reaction can be extended to the particular case of ZrCeTaSi2, where a solid solution grain is oxidized to the corresponding mixed oxy-carbide and then to the mixed oxide (5):  (Zr,Ta)C þ 2O2 / (Zr,Ta)CxOy þ C1\\x00x þ O4\\x00y / TaZr2.75O8 þ CO(5)  The formation of an intermediate oxy-carbide or cubic ZrO2 phase next to the cubic ZrC phase ensures that the scale adheres more firmly to the specimen substrate, owing to the same crystal structure, thereby improving its protective qualities. The partial  Fig. 8. SEM images of the surface of ZrCeTaSi2 after 20 min oxidation in air at (a) 1800, (b) 2000 and (c) 2200 K.  In the insets a magnification of the microstructure.  \\x0c', '414  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 9. SEM images of the polished cross-section of ZrCeTaSi2 after oxidation at 1800 K. (b), (c) Magnification of the areas as indicated in (a).  Fig. 10. SEM images of the fractured cross-section of ZrCeTaSi2 after oxidation at 2000 K. Only the outermost scale is shown.  sintering of the scale in the presence of carbon also impedes the diffusion of the reaction components and hence lowers the rate of oxidation [27]. The generated oxide layer is partially protective, as testified by the parabolic rate of gases evolution which is believed to control the oxide growth up to 1800 K. A different behaviour is observed upon oxidation at higher temperatures. Whilst for other carbides, like HfC the evolution of gases becomes less important owing to the sintering of the oxide [16,17], in the present case the gases evolution notably increases, probably owing to the melting of the mixed oxide. These considerations imply that the sintering and melting of the surface oxides determines a change in the oxidation behaviour. In particular, some authors recognized that the melting point of the TaZr2.75O8 phase could be significantly lower than that of pure zirconia with heavy consequences on the high temperature stability [28]. This assumption seems to find a confirmation in the test performed at 2200 K, where evident melting, dissociation, evaporation and re-precipitation phenomena occurred. The addition of significant amounts of silicides could partially alter the oxidation behaviour, so the oxidation reactions involving the Zr-silicides and the other secondary phases, like SiC, should be considered as well. Since the tests were performed from 1800 K on,  we can say that the surface of the specimens is mainly subjected to active oxidation regime that involves formation of gaseous products according to:  ZrSi2 þ 2O2 / ZrO2 þ 2SiO  SiC þ O2 / SiO þ CO  (6)  (7)  Gases formation introduces porosity which allows diffusion via pores. Gas formed below the oxide layer can also lift and disrupt the oxide layer, like in the present case from 2000 K on. In principle, oxidation of silicides should produce glassy silica that diffuses through the surface and form a stable and continuous silica layer. However, previous studies on the oxidation of similar composites have shown that even at temperatures lower than 1500 K, no continuous oxidation layer was observed on the surface, although silica partially filled cavities in the oxide cross-section [11]. This can be due to several reasons: there is not enough silica to fill all the volume expansion associated with the formation of the porous oxide layer and large CO escape resulting from oxidation of the carbide can further accelerate the dissociation of silica to gaseous SiO. Although no continuous silica layer was found on the  Fig. 11. SEM images of the polished cross-section of ZrCeTaSi2 after oxidation at 2200 K. (b)e(e) Magnification of the areas as indicated in (a) or (c).  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  415  surface, the presence of partially filled porosity in the cross-section can hinder the fast diffusion of gaseous species towards the unreacted bulk up to 2000 K (Fig. 10).  4.  Conclusions  A zirconium carbide ceramic was hot pressed at 1970 K with addition of tantalum silicide which enabled the complete densification. The matrix was composed by an inner ZrC-core surrounded by (Zr,Ta)C solid solution forming grains of about 5 mm. These were bordered by ZreSi phases, formed upon cation exchange between ZrC and TaSi2, and SiC particles, formed after carbo-reduction of the silica-based species. This ceramic displayed good mechanical properties with hardness of 18 GPa, fracture toughness of 3.6 MPa m1/2 and flexural strength of 500 MPa. Oxidation tests from 1800 to 2200 K for 20 min duration evidenced the formation of ZrO2 and of the mixed oxide TaZr2.75O8 with platelet shape. The ZrC-based composite resisted well up to 1800 K and until the first 10 min at 2000 K. After this controlled oxidation regime, concurrent phenomena of silicide decomposition, vigorous gas evolution and oxide melting took place, resulting in final oxide breaking and spalling.  Acknowledgements  We greatly acknowledge the financial support of the US Air Force Research Laboratory to partial of this activity through grant N. FA8655-12-1-3004, with Dr. Ali Sayir as contract monitor. D. Dalle Fabbriche is acknowledged for hot pressing, G. Celotti for X-ray diffraction and C. Melandri for mechanical testing.  References  [4]  [2]  [1] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, William Andrew Publishing/Noyes, 2001, pp. 55e99. Floyd B. Baker, Edmund K. Storms, Holley Jr., E. Charles, Enthalpy of formation of zirconium carbide, J. Chem. Eng. Data 14 (1969) 244e246. [3] H.F. Jackson, D.D. Jayaseelan, W.E. Lee, M.J. Reece, F. Inam, D. Manara, C. Perinetti Casoni, F. De Bruycker, K. Boboridis, Int. J. Appl. Ceram. Technol. 7 (2010) 316e326. B. Pierrat, M. Balat-Pichelin, L. Silvestroni, D. Sciti, High temperature oxidation of ZrCe20%MoSi2 in air for future solar receivers, Sol. Energy Mater. Sol. Cells 95 (2011) 2228e2237. E. Sani, L. Mercatelli, P. Sansoni, L. Silvestroni, D. Sciti, Spectrally selective ultra-high temperature ceramic absorbers for high-temperature solar plants, J. 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Robbins, H.F. McMurdie, Phase Diagram for Ceramists, American Ceramic Soc., Inc, Columbus, OH, 1964.  [21]  the  [9]  [23]  [24]  [25]  [26]  \\x0c']"
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  "_id": 287,
  "PDF": "Zirconium carbide doped with tantalum silicide- Microstructure, mechanical properties and high temperature oxidation.pdf",
  "Text": "['Materials Chemistry and Physics 143 (2013) 407e415  Contents lists available at ScienceDirect  Materials Chemistry and Physics  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t c h e m p h y s  Zirconium carbide doped with tantalum silicide: Microstructure, mechanical properties and high temperature oxidation  Laura Silvestroni a, *, Diletta Sciti a, Marianne Balat-Pichelin b, Ludovic Charpentier b  a ISTEC-CNR, Via Granarolo 64, 48018 Faenza, Italy b PROMES-CNRS, 7 rue du four solaire, 66120 Font-Romeu-Odeillo, France  h i g h l  i g h t s  \\x0f Densification of ZrC at 1970 K with addition of TaSi2. \\x0f Thorough study of the microstructure by TEM and hypothesis on the densification mechanism. \\x0f Mechanical characterization: HV e 18 GPa, KIc e 3.6 MPa m1/2, s e 500 MPa. \\x0f Oxidation tests at 1800e2200 K for 20 min and discussion on the various oxidation mechanisms.  a r t i c l e  i n f o  a b s t r a c t  A zirconium carbide ceramic was hot pressed to full density thanks to the addition of TaSi2, which enabled the densification to occur at 1970 K and improved the mechanical properties as compared to monolithic ZrC. The microstructure was analysed by combined X-ray diffraction, scanning and transmission electron microscopy to investigate the effective role of the sintering additive. In addition, high temperature oxidation was performed using the reactor REHPTS (Réacteur Hautes Pression et Température Solaire) from 1800 to 2200 K for 20 min and this composite demonstrated to resist towards the highly oxidative conditions better than other carbides, thanks to the chemical modification of the oxide formed upon Ta addition. However from 2000 K, the specimen resulted very damaged. Ó 2013 Elsevier B.V. All rights reserved.  Article history:  Received 11 January 2013 Received in revised form 12 September 2013 Accepted 18 September 2013  Keywords:  Carbides Electron microscopy Microstructure Mechanical properties Oxidation  1.  Introduction  Zirconium carbide (ZrC) is an extremely hard and refractory ceramic, belonging to the class of Ultra-High Temperature Ceramics (UHTCs) in view of its high melting point exceeding 3800 K [1]. Like all the components of this family, ZrC possesses a combination of interesting properties, like for example high corrosion resistance in both acid and basic environments, a low thermal conductivity \\x001 K \\x001) and high electrical conductivity thanks to the (20.5 W m presence of metallic bonds. Moreover, the strong covalent ZreC bonds confer to ZrC high hardness (25 GPa), high Young’s modulus (440 GPa) and mechanical strength. Being a Zr-based compound, it possesses lower density compared to other transition metal carbides, like WC, TaC or HfC. Similar to other transition metal carbides, ZrC is often sub-stoichiometric and has a stability range of  * Corresponding author. laura.silvestroni@istec.cnr.it (L. Silvestroni).  E-mail address:  0254-0584/$ e see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matchemphys.2013.09.020  carbon to metal ratio from 0.65 to 0.98; when carbon content exceeds 0.98, the material contains free carbon too [2]. It is commercially used in tool bits for cutting and it is potentially suitable for re-entry vehicles, rockets or scramjet engines, where low density and high temperature load bearing capability are required. Other applications of ZrC ceramics can be found in nuclear sector, where it is used as refractory coating in reactors, thanks to its low neutron absorption cross-section and weak damage sensitivity under irradiation [3]. It has also been recognized that ZrC possesses favourable emissivity properties rendering it a promising material for use as absorber in concentrating solar power systems [4e6]. Besides these amazing chemico-physical properties, the two main drawbacks of ZrC consist in a difficult densification and a poor oxidation resistance, which characterizes all the transition metal carbides. Sintering is generally performed by pressure assisted techniques, like hot pressing (HP) [7] or spark plasma sintering (SPS) at very high temperatures [3,8,9], inducing however coarse microstructure with grain size in the order of 10e30 mm. Other  \\x0c', '408  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  pressureless sintering technologies enable the densification of ZrC at lower temperatures, but only upon addition of sintering additives such as MoSi2, ZrB2, HfC, carbon or SiC [10,11]. On the other side, poor oxidation resistance over 1070 K strongly limits the applications of ZrC and makes it necessary to work in protective environments. The final oxidation products of ZrC are ZrO2 and carbon oxides. ZrO2 forms a fine grained porous scale which allows gaseous diffusion of O2 through the pores to the ZrC surface, and therefore provides no oxidation protection. In addition, low-temperature formation of cubic modification of ZrO2, instead of thermodynamically stable monoclinic modification, may significantly contribute to the high oxidation rate as a result of about 5 orders of magnitude higher solid state oxygen diffusion in cubic zirconia compared to that in monoclinic [12]. One way to improve the oxidation resistance of ZrC is to make composites, adding for example ZrB2, SiC or silicides, able to hinder oxygen penetration through the B2O3-filled ZrO2 layer, or to form a protective silica scale [13]. In this work, TaSi2 was chosen as sintering additive for ZrC to enable densification at temperature lower than 2370 K and improve the oxidation resistance of the carbide matrix. The selection of TaSi2 was driven by its refractoriness, as it melts at 2370 K, and its proved efficacy as sintering additive for other transition metal carbides [14]. In addition, previous works on ZrB2 enabled to conclude that Ta-addition is beneficial for oxidation up to 1870 K owing to its capability to stuff oxygen vacancies in Zroxide by a higher valence cation, thus rendering the oxide layer more stable in the tetragonal crystal structure [15], or thanks to the tendency to increase the glass viscosity and induce immiscibility, which is translated in higher boiling point of the glass [16]. However tantalum becomes detrimental above 2150 K due to the melting of its oxide, Ta2O5 [17], or to the formation of a complex orthorhombic oxide, TaZr2.75O8, which has a needle-like morphology not favourable for adhesion to the unreacted bulk [15,18] and a lower melting point than pure tantalumor pure zirconium-oxides [19]. In this work, the room temperature mechanical properties of the dense ZrCeTaSi2 were measured and compared to other ZrCcompounds, then the oxidation behaviour of the ceramic was studied in the temperature range from 1800 to 2200 K in order to evaluate the performances at the different temperatures and identify eventual operating limits.  2.  Experimental procedure  2.1. Material  The ceramic was produced starting from the following commercial powders in amount 85 vol% ZrC and 15 vol% TaSi2:  cubic ZrC (H. C. Starck, Germany): Grade B, mean particle size: 3.5 mm; particle size range 0.8e6 mm; specific surface area 3 m2 g \\x001; impurities (wt%): C 1.5, O 0.6, N 0.8, Fe 0.05, Hf 2.0. hexagonal TaSi2 (Cerac Inc., Milwaukee, WI), purity 99.5%, \\x00325 mesh, particle size range 5e10 mm.  The powder mixture was ball milled for 24 h in absolute ethanol using silicon carbide milling media. Subsequently the powders were dried in a rotary evaporator, sieved through a 60-mesh screen and green shaped to 45 mm-pellet by linear pressing at 20 MPa. Hot-pressing was conducted in low vacuum (w100 Pa) using an induction-heated graphite die with a constant uniaxial pressure of 30 MPa, maximum sintering temperature of 1970 K, heating rate of \\x001 and free cooling. 20 K min  The final density was measured by Archimedes’ method on a hydrostatic balance. The theoretical density was calculated by the rule of mixtures. Crystalline phases were identified by X-ray diffraction (Siemens D500, Germany) using the CuKa radiation, with 0.04 2q-step size and 1 s scan step time. The microstructure was analysed using scanning electron microscopy (SEM, Cambridge S360) and energy dispersive spectroscopy (EDS, INCA Energy 300, Oxford Instruments, UK) on fractured and polished surfaces. TEM samples were prepared by cutting 3 mm discs from the sintered pellets. These were mechanically ground down to about 20 mm and then further ion beam thinned until small perforation were observed by optical microscope. The detailed phase analysis was performed using a transmission electron microscopy (FEI Tecnai F20 ST), with an acceleration voltage of 200 kV, equipped with an EDAX EDS X-ray spectrometer PV9761 with Super ultra-thin window. Microstructural parameters, like amount of residual porosity or secondary phases, were determined through image analysis on the micrographs of polished surfaces by commercial software (Image Pro-plus 4.0, Media Cybernetics, Silver Springs, USA).  2.2. Mechanical properties  Vickers microhardness (HV) was measured on the polished surface, with a load of 9.81 N, using a Zwick 3212 tester. Fracture toughness (KIc) was evaluated using chevron-notched 25 \\x02 2 \\x02 2.5 mm beam (CNB) in flexure. The test bars, (length \\x02 width \\x02 thickness, respectively), were notched with a 0.08 mm diamond saw; the chevron-notch tip depth and average side length were about 0.12 and 0.80 of the bar thickness, respectively. The flexural tests were performed on a semi-articulated silicon carbide four-point jig with a lower span of 20 mm and an upper span of 10 mm on a universal screw-type testing machine (Instron mod. 6025). The specimens were deformed with a cross\\x001. The slice model equation of Munz head speed of 0.05 mm min et al. was used to calculate KIc [20]. On the same machine and with the same flexural jig, the flexural strength (s), was measured on chamfered bars 25 \\x02 2.5 \\x02 2 mm (length \\x02 width \\x02 thickness, respectively), using a crosshead speed \\x001; five specimens were tested. of 0.5 mm min  2.3. High temperature oxidation resistance  Pression  et  Température  Solaire, High  (Réacteur Hautes  The reactor used to perform high temperature oxidation is the REHPTS Pressure and Temperature Solar Reactor), implemented at the focus of the Odeillo 5 kW solar furnace [21]. A heliostat reflects the incident solar flux to a concentrator with faceted mirrors. A shutter enables to control the fraction of the concentrated solar flux delivered to the sample placed inside the reactor and therefore its surface temperature. In this set-up the disc (Ø ¼ 25 mm, h ¼ 2 mm) is placed 25 mm above the focus of the solar furnace, so that very high temperatures may be achieved on a homogeneous 10 mm diameter area and at very fast rate, up to \\x001. Two mirrors enable a monochromatic (5 mm) optical 100 K s pyrometer (Ircon, Modline Plus) to measure the surface temperature of the sample through a fluorine window on a 6 mm diameter circle on the sample surface at a pyrometer-sample distance of 120 cm. To obtain the real temperature by monochromatic pyrometry, a normal spectral emissivity of 0.75 was used in these experiments, as the sample was immediately covered by an oxide layer mainly composed of zirconia. The pyrometer, together with all the parts present on the optical path, was calibrated on a blackbody.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  409  is  from  going  The accuracy of the temperature measurement 1400 \\x06 15 K to 2200 \\x06 22 K. The oxidations were performed in air with an atmosphere continuously renewed. Due to the altitude of the laboratory, the total atmospheric pressure is around 87 kPa and the oxygen partial pressure pO2 is 17 kPa. The temperature of the samples was maintained at a constant plateau during 20 min and a video camera was used to follow in situ the oxidation process. A mass spectrometer (Pfeiffer Omnistar) enabled in situ gas phase analysis. CO is expected to be one of the main gaseous products during oxidation, but its molar weight is the same as N2 one (m/e ¼ 28), so it is impossible to separate the contribution of CO from the one of preponderant N2. We therefore mainly followed the signal corresponding to m/e ¼ 44, corresponding both to CO2 and gaseous SiO. The samples were weighted before and after oxidation in order to assess the mass variation and the surfaces and cross-sections were analysed after oxidation using XRD, SEM and EDS.  3. Results and discussion  3.1. As sintered microstructure  The addition of TaSi2 to ZrC matrix enabled the achievement of \\x003, at 1970 K, which is a notably lower temfull density, 7.11 g cm perature as compared to conventional ZrC-based ceramics for which 2170e2370 K are required, see Table 1. X-ray diffraction did not evidence the presence of either ZrC or TaSi2. On the contrary a new phase with reduced cell parameter was identified as a (Zr,Ta)C solid solution (Fig. 1). The substitution of Zr atoms by Ta from TaSi2 provoked a shrinkage of the lattice from 4.692 \\x17A for ZrC to 4.646 \\x17A, which corresponds to a carbide with formula (Zr0.8Ta0.2)C, according to Vegard’s rule. The final microstructure of this composite was very complex, as the SEM images display in Fig. 2. Fig. 2a shows a homogeneous and dense microstructure with around 5 vol% of black particles of SiC. In Fig. 2b a magnification of the polished surface also shows dark grey phases with irregular faceted shape containing ZreTae SieC in different amounts, probably deriving from the original TaSi2, which is instead not clearly detected. The inset of Fig. 2b evidences the morphology of the matrix grain. As a general guideline, bright contrast indicates higher Ta content, whilst darker contrast indicates higher Zr percentage. The core with jagged edges is ZrC, moving outwards the contrast becomes brighter and EDS revealed the presence of Ta, giving rise to an approximate formula (Zr0.85Ta0.15)C. Inside the grain, 500 nm particles of ZreTaeSieC phases are visible, but the stoichiometry is difficult to define. The mean grain size of the newly formed carbide is around 5.2 mm with grains achieving up to 8 mm. Despite  the mean  being sintered at relatively low temperatures, 1970 K, grain size was notably coarsened. TEM analysis allowed a deeper investigation of the matrix: Fig. 3a evidences the coreeshell morphology with denticulate interface; the EDS spectra in Fig. 3b show the chemical composition of the two regions. Particles incorporated in SiC large grains were identified as Zr-rich silicides, Zr2Si or Zr5Si3, as those displayed in Fig. 3c. At the triple points, further ZreSi phases with flat edges were recognized, as the examples reported in Fig. 3d,e. High resolution imaging revealed the presence of dislocations between ZrC-core and (Zr,Ta)C-shell (Fig. 4a) and generally clean grain boundaries between adjacent (Zr,Ta)C grains or in contact with Zrsilicides (Fig. 4b,c) were found. Occasionally, amorphous SiO2 was noticed at the grain boundaries between the carbide and the silicide phases (Fig. 4d).  3.2. Densification behaviour  With the aim to identify the densification mechanisms activated by TaSi2, the main outcomes regarding the microstructural evolution of ZrCeTaSi2 can be summarized as follows:  ZrC original grains were surrounded by Ta-containing solid solutions, which grew epitaxially on the matrix grain (Fig. 3a). The misfit between ZrC and the (Zr,Ta)C solid solution was accommodated by corrugated grain boundaries and dislocations (Fig. 4a). These features suggest great solubility of Ta in ZrC; after sintering, no TaSi2 was observed anymore, but ZreSi phases with low dihedral angles took its place, revealing a mutual solubility of Zr in TaeSi-based phases. The new formed silicides were located mainly at the triple points (Fig. 3e). High resolution TEM showed evidence of generally clean grain boundaries between carbide matrix and silicide (Fig. 4a).  Previous studies of the interaction between carbides and TaSi2 revealed that during sintering, in presence of CO, TaSi2 tends to dissociate and decompose into Ta, gaseous SiO, liquid Si and eventually to form new TaC phase [14], according to reaction (1). The formation of liquid silicon at relatively low temperature is compatible with the enhanced sintering activity of this composite which started to shrink below 1770 K, close to the melting temperature of Si.  TaSi2 þ CO(g) / TaC þ SiO(g) þ Si(l)  (1)  Free Si was not clearly observed in the final microstructure; however in the highly reducing environment of the hot pressing chamber with graphite dies and rams, it is very probable that liquid silicon was reduced to SiC, as testified by Fig. 2:  Table 1 Sintering parameters, microstructural features and mechanical properties of some ZrC-based ceramics. Legend: HP-hot pressing, SPSspark plasma sintering, PLSepressureless sintering, m.g.s. e mean grain size, HV e hardness, KIceCNB fracture toughness, s e 4-pt flexural strength.  Label  Sintering  ZCT ZCM ZCM ZCM ZCZg ZCS  K, min, atm, MPa  HP 1970, 6, vac, 30 HP 2170, 12, vac, 30 SPS 1970, 3, vac, 100 PLS 2220, 60, Ar, HP 2170, 60, vac, 30 PLS 2370, 120, Ar,  a Direct crack measurements. b 3-pt bending.  Sint. add.  vol%  15 TaSi2 15 MoSi2 9 MoSi2 20 MoSi2 8 wt Zr þ 2.3 wt graphite 20 SiC  Rel. r  %  99.9 96.8 99.0 96.8 98.3 96.7  m.g.s  mm  5.2 3.9 3.5 6.0 w10 3.1  HV  GPa 17.9 \\x06 0.7 19.2 \\x06 0.4 20.0 \\x06 0.5 12.7 \\x06 1.0 16.2 \\x06 0.9 11.8 \\x06 0.8  KIc  (MPa m1/2) 3.6 \\x06 0.4 3.4 \\x06 0.5 3.3 \\x06 0.4a 3.5 \\x06 0.2 4.7 \\x06 0.4a  e  s  MPa 503 \\x06 55 474 \\x06 41 591 \\x06 48b 272 \\x06 12  e  474b  Ref  This work [23] [8] [11] [7] [25]  \\x0c', '410  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 1. X-ray diffraction pattern of the as sintered ZrCeTaSi2 composite showing peaks shift from the pure ZrC phase.  2Si(l) þ CO(g) / SiC þ SiO(g)  (2)  Those findings are in good accordance with the microstructural features, as pure TaSi2 is not found anymore and TaC is instead present in the solid solution with ZrC according to Eq. (3):  The well-defined boundary between core and shell and the morphology of the interface between them with trapped particles also put forward a re-precipitation from liquid phase over a diffusion process.  TaC þ ZrC / (Zr,Ta)C  3.3. Mechanical properties  (3)  Indeed, it is well known that the solubility between carbides of Group IV and mono-carbides of Group V is complete and they are expected to form solid solutions [1]. Microstructural features and thermodynamics suggest that densification occurred through transient liquid phase. This hypothesis is further strengthened by the irregular shape displayed by the residual silicide phases, as well as the relevant coarsening of the carbide grains. We can reasonably suppose that liquid phase containing TaeSieCeO, where Zr was soluble in, formed during densification and crystallized at the triple points in form of ZreSi phases upon cooling leaving clean grain boundaries. As solid solutions formed, Ta from TaSi2 substituted Zr atoms in ZrC lattice. This may occur either by cations diffusion or by solution-reprecipitation. Given the low self-diffusion coefficient of this class of materials, it is presumed that lattice diffusion can occur only at very high temperature. Indeed, solution re-precipitation seems to be the dominant mechanism, in light of the sintering behaviour characterized by a relatively low shrinkage temperature.  Table 1 summarizes the main sintering parameters, microstructural features and mechanical properties of the ZrCeTaSi2 ceramic, compared to other ZrC-based materials taken from literature. The hardness of the ZrCeTaSi2 composite was about 18 GPa, in the range of the data reported in the literature for similar materials. The discrepancy between monolithic ZrC, 25 GPa, and this composite is due to the presence of softer silicide phases which have hardness below 16 GPa [22] and the quite coarse microstructure. The fracture toughness resulted in the range or higher than the other ZrC-composites, indicating that this property is mainly dictated by ZrC matrix toughness [7,8,11,23,24]. ZrCeTaSi2 composite displayed strength around 500 MPa, higher than similar composites containing MoSi2 processed by hot pressing and pressureless sintering [7,8,11,23,24]. The highest value shown in Table 1 for the SPS-composite refers to measurements carried out on 3-point bending and very small specimens [8]. So it is concluded that this ceramic displayed very good mechanical properties also after sintering at such low temperature (1970 K).  Fig. 2. SEM images of the polished surface of ZrCeTaSi2 composite showing (a) the formation of SiC dark particles and (b) the morphology of the matrix.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  411  Fig. 3. TEM images of the as sintered ZrCeTaSi2 composite showing (a) the morphology of the matrix grain with the corresponding EDS spectra in (b), (c) SiC grains incorporating ZreSi phases, (d) and (e) ZreSi phases at the triple point.  3.4. Microstructure of oxidized specimens  Pictures of the samples after testing at 1800, 2000 and 2200 K for 20 min, are shown in Fig. 5. The front faces are covered with a grey layer becoming progressively white indicating that the expected oxidation occurred. The oxide layer that formed at 1800 K, Fig. 5a, looks quite smooth and well adherent to the carbide, indicating that at this stage the sample well survived the test. At  2000 K, Fig. 5b, the sample broke upon cooling and some bubbles can be seen in the more heated zone. At 2200 K, Fig. 5c, bubbling phenomena are more evident, the surface is whiter and a wavy surface is also noticeable, indicating a change in the oxidation resistance at this temperature. Fig. 6 shows the evolution of CO2 and SiO as a function of the oxidation temperature, determined using mass spectrometry. We can observe that at 1800 K the material evolves very low amount of  Fig. 4. HR-TEM images of the as sintered ZrCeTaSi2 composite showing (a) the interface between ZrC-core and (Zr,Ta)C shell with dislocations indicated by arrows, (b), (c) examples of clean grain boundaries between (Zr,Ta)C e (Zr,Ta)C e ZrSi2, or partially wetted grain boundaries in (d).  \\x0c', '412  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 5. Photographic images of the ZrCeTaSi2 discs after oxidation tests at (a) 1800, (b) 2000 and (c) 2200 K.  Fig. 9 presents the cross-section of the sample oxidized at 1800 K, where a 140 mm thick layer underwent oxidation with formation of the mixed Zr, Ta oxide with granular shape covered by 3 mm thin and discontinuous silica layer (Fig. 9b). Moving further inward, down to around 400 mm from the surface, a complex mixture of (Zr,Ta)Si2, (Zr,Ta)-oxy-carbide and SiC phases were found, Fig. 9c. The complex (Zr,Ta)eCeO phase is the result of partial oxidation of the starting matrix grain, made also of (Zr,Ta)C, as outlined above. The interfaces between the oxide, oxy-carbide and the bulk are crack free and continuous. The formation of an oxy-carbide standing between the pure oxide external layer and the carbide core, had been already reported for ZrC and HfCcompounds [25,26]. After oxidation at 2000 K, the specimen resulted in a layered structure with the outermost scale composed of 5 mm long TaZr2.75O8 grains standing out on a continuous silica-based scale (Fig. 10b), which topped about 300 mm of coarse ZrO2 and TaZr2.75O8 grains, partially filled with silica and where vigorous bubbling clearly occurred Fig. 10a. This thick external layer was partially detached from the underneath surface, which is not shown, since it had the same aspect as the section oxidized at 1800 K (Fig. 9). The cross-section of the ceramic oxidized at 2200 K is shown in Fig. 11, where all the thickness of the disc resulted partially oxidized. The outermost thick scale is composed by a compact ZrO2 layer where 20e30 mm large porosities can be found (Fig. 11a). In this region, magnified in Fig. 11b, SiO2 droplets (dark) containing Zr traces are surrounded by a brighter phase, identified as a solid  gaseous species, which notably increases after 10 min of oxidation at 2000 K. This trade off can be due to the migration of silicide phases to the surface, which form a protective glassy scale until the dissociation of ZreSi silicides prevails, resulting in further development of SiO. At 2200 K the gases escaping is notably higher, owing to the complete dissociation of the silica-glass and the lack of a protective oxide scale, or to CO2 escape through fissures. As far as concerns mass variation, as expected, the weight gain increased with increasing temperature going from \\x002 min \\x001 at 1800 K, to 2.37 mg cm \\x002 min \\x001 at 2000 K 1.20 mg cm \\x002 min \\x001 at 2200 K. and 3.85 mg cm X-ray diffraction patterns collected on the surface of the three specimens after oxidation tests are shown in Fig. 7. The main crystalline phases at all temperatures is a mixed oxide with composition TaZr2.75O8, which has an orthorhombic structure and preferred orientation along the [020] planes; the peaks become sharper increasing the temperature, indicating improved crystallization. Monoclinic ZrO2 is also present and the peaks signal increases with temperature to about 25%1 at 2200 K, indicating the higher stability of pure oxide over the mixed one. Cubic and tetragonal ZrO2 could be present too, but the superimposition of the main peaks with the mixed oxide hinders a conclusive analysis. Other authors [25,26] and similar tests on ZrC-based composites showed indeed that the presence of carbon, coming from the oxidation of the carbide, stabilizes c-ZrO2 at low temperatures. Fig. 8 presents SEM images of the surfaces of the central region of the discs at 1800, 2000 and 2200 K. The addition of TaSi2 to a ZrC matrix generated a variety of elaborated morphologies varying the oxidation temperature; as a rule of thumb, dark regions correspond to silica-based glass, bright phases to oxides. At 1800 K the surface of the composite presents a rough cracked aspect with TaZr2.75O8 and ZrO2 being the main phases where discrete pockets of silica-based glass are found, Fig. 8a. The surface at 2000 K is mainly composed of petal-like grains of TaZr2.75O8 which form volcanos and tend to microcracking, Fig. 8b. At 2200 K melting and recrystallization of the oxide occurred and 20 mm large grains formed leaving residual silica, containing Zr traces, at the grain boundaries, Fig. 8c. These grains are in turn composed of polyhedral structures of ZrO2 and TaZr2.75O8 showing the growth planes decorated by a dendritic irregularly shaped phase identified as ZrO2 and ZrSiO4 containing small traces of Ta, inset in Fig. 8c.  1  Since the scattering coefficient for the orthorombic TaZr2.75O8 phase has not been published in the literature and is not available in the ICSD database, we considered this phase as a solid solution between 0.5 mol orthorombic Ta2O5 (#54e 514) and 3 mol tetragonal ZrO2 (#42.1164) and estimated a scattering coefficient of 4.7.  Fig. 6. Measured concentration of SiO and CO2 (m/e ¼ 44) produced at (a) 1800, (b) 2000 and (c) 2200 K during the oxidation of ZCT.  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  413  Fig. 7. X-ray diffraction patterns of the ZrCeTaSi2 composite after oxidation at (a) 1800, (b) 2000 and (c) 2200 K.  solution with possible formula (Zr0.8Ta0.2)O2. Right underneath this 550 mm layer, the mixed TaZr2.75O8 oxide is present in form of a thick dense scale including elongated porosity and a Ta-rich white oxide (Fig. 11c, upper part). Moving further inward, fine grained ZrO2 with Ta traces and SiO2 discrete phases (Fig. 11d) stand above the already mentioned (Zr,Ta)-oxy-carbide and SiC phases (Fig. 11c lower part and Fig. 11e). The core of the disc is composed by a mixed (Zr,Ta)-oxy-carbide containing progressively lower oxygen amounts. These articulate morphologies are the result of complex oxidation mechanisms, including melting, phase separation and reprecipitation, occurring at 2200 K.  3.5. Oxidation behaviour  Pictures and SEM images clearly reveal this composite to undergo various oxidation mechanisms depending on the different temperature range. Fig. 6 evidences that the evolution of CO2 and SiO gaseous species has a slow and parabolic trend at 1800 K for the whole duration of the test, but at 2000 K an abrupt increase occurs after 10 min of oxidation. This variation was ascribed to the melting and decomposition of Zr-silicides, occurring at around 1900 K [27], which resulted in further SiO release. A second explanation could  be the migration of silica glass to the surface, as observable in Fig. 10b, which after a certain period loses its shielding action and is no longer protective to gases escape, as indicated by the large voids in the cross-section of Fig. 10a, and by the bubbling and fracture of the sample in Fig. 5b. From this temperature on, vigorous gas escape was reported to occur, as demonstrated also by the turbulent microstructure in Fig. 11a. It is almost established that the oxidation of carbides of groups IVeVI transition metals occurs through formation of an oxy-carbide of the metal plus carbon, which is subsequently oxidized to CO and CO2, ending with formation of the metal oxide [25,26], according to reactions of the type:  MeC þ 2O2 / MeCxOy þ C1\\x00x þ O4\\x00y / MeO2 þ CO2  (4)  This general reaction can be extended to the particular case of ZrCeTaSi2, where a solid solution grain is oxidized to the corresponding mixed oxy-carbide and then to the mixed oxide (5):  (Zr,Ta)C þ 2O2 / (Zr,Ta)CxOy þ C1\\x00x þ O4\\x00y / TaZr2.75O8 þ CO(5)  The formation of an intermediate oxy-carbide or cubic ZrO2 phase next to the cubic ZrC phase ensures that the scale adheres more firmly to the specimen substrate, owing to the same crystal structure, thereby improving its protective qualities. The partial  Fig. 8. SEM images of the surface of ZrCeTaSi2 after 20 min oxidation in air at (a) 1800, (b) 2000 and (c) 2200 K.  In the insets a magnification of the microstructure.  \\x0c', '414  L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  Fig. 9. SEM images of the polished cross-section of ZrCeTaSi2 after oxidation at 1800 K. (b), (c) Magnification of the areas as indicated in (a).  Fig. 10. SEM images of the fractured cross-section of ZrCeTaSi2 after oxidation at 2000 K. Only the outermost scale is shown.  sintering of the scale in the presence of carbon also impedes the diffusion of the reaction components and hence lowers the rate of oxidation [27]. The generated oxide layer is partially protective, as testified by the parabolic rate of gases evolution which is believed to control the oxide growth up to 1800 K. A different behaviour is observed upon oxidation at higher temperatures. Whilst for other carbides, like HfC the evolution of gases becomes less important owing to the sintering of the oxide [16,17], in the present case the gases evolution notably increases, probably owing to the melting of the mixed oxide. These considerations imply that the sintering and melting of the surface oxides determines a change in the oxidation behaviour. In particular, some authors recognized that the melting point of the TaZr2.75O8 phase could be significantly lower than that of pure zirconia with heavy consequences on the high temperature stability [28]. This assumption seems to find a confirmation in the test performed at 2200 K, where evident melting, dissociation, evaporation and re-precipitation phenomena occurred. The addition of significant amounts of silicides could partially alter the oxidation behaviour, so the oxidation reactions involving the Zr-silicides and the other secondary phases, like SiC, should be considered as well. Since the tests were performed from 1800 K on,  we can say that the surface of the specimens is mainly subjected to active oxidation regime that involves formation of gaseous products according to:  ZrSi2 þ 2O2 / ZrO2 þ 2SiO  SiC þ O2 / SiO þ CO  (6)  (7)  Gases formation introduces porosity which allows diffusion via pores. Gas formed below the oxide layer can also lift and disrupt the oxide layer, like in the present case from 2000 K on. In principle, oxidation of silicides should produce glassy silica that diffuses through the surface and form a stable and continuous silica layer. However, previous studies on the oxidation of similar composites have shown that even at temperatures lower than 1500 K, no continuous oxidation layer was observed on the surface, although silica partially filled cavities in the oxide cross-section [11]. This can be due to several reasons: there is not enough silica to fill all the volume expansion associated with the formation of the porous oxide layer and large CO escape resulting from oxidation of the carbide can further accelerate the dissociation of silica to gaseous SiO. Although no continuous silica layer was found on the  Fig. 11. SEM images of the polished cross-section of ZrCeTaSi2 after oxidation at 2200 K. (b)e(e) Magnification of the areas as indicated in (a) or (c).  \\x0c', 'L. Silvestroni et al.  / Materials Chemistry and Physics 143 (2013) 407e415  415  surface, the presence of partially filled porosity in the cross-section can hinder the fast diffusion of gaseous species towards the unreacted bulk up to 2000 K (Fig. 10).  4.  Conclusions  A zirconium carbide ceramic was hot pressed at 1970 K with addition of tantalum silicide which enabled the complete densification. The matrix was composed by an inner ZrC-core surrounded by (Zr,Ta)C solid solution forming grains of about 5 mm. These were bordered by ZreSi phases, formed upon cation exchange between ZrC and TaSi2, and SiC particles, formed after carbo-reduction of the silica-based species. This ceramic displayed good mechanical properties with hardness of 18 GPa, fracture toughness of 3.6 MPa m1/2 and flexural strength of 500 MPa. Oxidation tests from 1800 to 2200 K for 20 min duration evidenced the formation of ZrO2 and of the mixed oxide TaZr2.75O8 with platelet shape. The ZrC-based composite resisted well up to 1800 K and until the first 10 min at 2000 K. After this controlled oxidation regime, concurrent phenomena of silicide decomposition, vigorous gas evolution and oxide melting took place, resulting in final oxide breaking and spalling.  Acknowledgements  We greatly acknowledge the financial support of the US Air Force Research Laboratory to partial of this activity through grant N. FA8655-12-1-3004, with Dr. Ali Sayir as contract monitor. D. Dalle Fabbriche is acknowledged for hot pressing, G. Celotti for X-ray diffraction and C. Melandri for mechanical testing.  References  [4]  [2]  [1] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, William Andrew Publishing/Noyes, 2001, pp. 55e99. Floyd B. Baker, Edmund K. Storms, Holley Jr., E. Charles, Enthalpy of formation of zirconium carbide, J. Chem. Eng. Data 14 (1969) 244e246. [3] H.F. Jackson, D.D. Jayaseelan, W.E. Lee, M.J. Reece, F. Inam, D. Manara, C. Perinetti Casoni, F. De Bruycker, K. Boboridis, Int. J. Appl. Ceram. Technol. 7 (2010) 316e326. B. Pierrat, M. Balat-Pichelin, L. Silvestroni, D. Sciti, High temperature oxidation of ZrCe20%MoSi2 in air for future solar receivers, Sol. Energy Mater. Sol. Cells 95 (2011) 2228e2237. E. Sani, L. Mercatelli, P. Sansoni, L. Silvestroni, D. Sciti, Spectrally selective ultra-high temperature ceramic absorbers for high-temperature solar plants, J. 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Alper, High Temperature Oxides Part II Oxide of Rare Earths, Titanium, Zirconium, Hafnium, Niobium and Tantalum, Academic Press, New York and London, 1970. E. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2 and HfB2-based ultra-high temperature ceramics: Effect of Ta additions, J. Mater. Sci. 39 (19) (2004) 5969e5977. [19] A.K. Bhattacharya, V. Shklover, W. Steurer, G. Witz, H.P. Bossmann, O. Fabrichnaya, Ta2O5eY2O3eZrO2 system: experimental study and preliminary thermodynamic description, J. Eur. Ceram. Soc. 31 (3) (2011) 249e 257. [20] D.G. Munz, J.L. Shannon Jr., R.T. Bubsey, Fracture Toughness Calculation From Maximum Load in Four Point Bend Tests of Chevron Notch Specimens, Int. J. Fract 16 (3) (1980) R137e141. L. Charpentier, K. Dawi, J. Eck, B. Pierrat, J.L. Sans, M. Balat-Pichelin, Concentrated solar energy to study high temperature materials for space and energy, J. Sol. Energy Eng. 133 (3) (2011) 031005 (8pp.). [22] G. Schultes, M. Schmitt, D. Goettel, O. Freitag-Weber, Strain sensitivity of TiB2, TiSi2, TaSi2 and WSi2 thin films as possible candidates for high temperature strain gauges, Sens. Actuators A 126 (2006) 287e291. L. Silvestroni, D. Sciti, Effect of transition metal silicides on microstructure and mechanical properties of ultra-high temperature ceramics, in: J. Low, Y. Sakka, C. Hu (Eds.), MAX Phases and Ultra-high Temperature Ceramics for Extreme Environments, IGI Global, Hershey, PA, 2013, pp. 125e179. L. Zhao, D. jia, X. Duan, Z. Yang, Y. Zhou, pressureless sintering of ZrC-based ceramics by enhancing power sinterability, Int. J. Refract. Met. Hard Mater. 29 (2011) 516e521. F. Voitovich, E.A. Pugach, High-temperature oxydation of ZrC and HfC, Powder Metall. 11 (131) (1973) 67e74. Shiro Shimada, Michio Inagaki, Mikio Suzuki, Microstructural observation of the ZrCyZrO2 interface formed by oxidation of ZrC, JMR 11 (10) (1996) 2594e 2597. [27] H. Okamoto, Bull. Alloy Phase Diagrams 11 (1990) 513e519. [28] E.M. Levin, C.R. Robbins, H.F. McMurdie, Phase Diagram for Ceramists, American Ceramic Soc., Inc, Columbus, OH, 1964.  [21]  the  [9]  [23]  [24]  [25]  [26]  \\x0c']"
},{
  "_id": 288,
  "PDF": "Zirconium carbide oxidation - Kinetics and oxygen diffusion through the intermediate layer.pdf",
  "Text": "['O R I G I N A L A R T I C L E  Zirconium carbide oxidation: Kinetics and oxygen diffusion through the intermediate layer  Claudia Gasparrini1  | Richard J. Chater2  | Denis Horlait1,3  |  Luc Vandeperre2  |  William E. Lee1  1Centre for Nuclear Engineering (CNE)  & Department of Materials,  Imperial  College London, London, UK  2Department of Materials,  Imperial  College London, London, UK  3UMR 5797, CNRS, Centre D’Etudes Nucl\\x13eaires de Bordeaux-Gradignan,  Gradignan, France  Correspondence  Claudia Gasparrini, Centre for Nuclear  Engineering (CNE) & Department of  Materials,  Imperial College London,  London, UK.  Email: cg1614@ic.ac.uk  Funding information  Engineering and Physical Sciences  Research Council, Grant/Award Number:  DISTINCTIVE EP/L014041/1,  Industrial  Case Award EP/M507428/1  Abstract  Oxidation of hot-pressed ZrC was  investigated in air  in the 1073-1373 K range.  The kinetics were  linear  at 1073 K, whereas  at higher  temperature  samples  ini tially followed linear kinetics before undergoing rapid oxidation leading to a Mal tese  cross  shape of  the oxide. The  linear kinetics  at 1073 K was governed by  inward  oxygen  diffusion  through  an  intermediate  layer  of  constant  thickness  between ZrC and ZrO2 which was  comprised  of  amorphous  carbon  and ZrO2  nanocrystals. Diffusion of oxygen through the intermediate layer was measured to \\x001 using 18O as a tracer  be 9 9 10  \\x0010 cm2 s  in a double oxidation experiment  in  16O/18O. Oxidation at 1073 and 1173 K produced samples made of m-ZrO2 and  either  tor c-ZrO2 with an adherent  intermediate layer made of amorphous carbon  and ZrO2, whereas oxidation at 1273 and 1373 K produced samples with a volu minous oxide made of m-ZrO2 showing a gap between ZrC and the oxide. A sub stoichiometric  zirconia  layer was  found  at  the  gap  at  1273 K and  no  carbon  uptake was detected in this  layer when compared with the  top oxide  layer. The  loss  of  the intermediate layer and the slowdown of the linear rate constant \\x001) at 1273 K compared to 1173 K was correlated with the preferential  (g m  \\x002 s  oxidation of  carbon at  the  intermediate  layer which would leave  as CO and/or  CO2 leaving a gap between ZrC and substoichiometric zirconia.  K E Y W O R D S  oxidation, zirconia, zirconium/zirconium compounds  1  |  I N T R O D U C T I O N  Zirconium carbide  has potential in ultrahigh-temperature for coatings on hypersonic vehicles1,2 industry.3,4 ZrC  applications  such as  or  as  an inert matrix fuel  in the  nuclear  has better properties than other carbides such as a higher thermal conductivity and higher melting point,5-8 however,  a comprehensive database on its properties and therefore its use is rarely considered.9  is not available  The oxidation of ZrC has been studied10-19 under several experimental conditions of temperature and pressure1014 but a clear understanding has not yet been achieved due  to  the  large  numbers  of  variables  that  affect  it,  such  as  chemical composition of the initial material chiometry), porosity, or grain size.9,20 A mechanistic model recently proposed by Katoh et al9 building on the model previously proposed by Shimada21  (C/Zr  stoi showed as  an initial  stage  the  formation of  an oxycarbide   This is an open access article under  the terms of  the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the  original work is properly cited. © 2018 The Authors. Journal of  the American Ceramic Society published by Wiley Periodicals,  Inc. on behalf of American Ceramic Society (ACERS)  Received: 28 November 2017  |  Accepted: 1 February 2018  DOI: 10.1111/jace.15479  2638  |  wileyonlinelibrary.com/journal/jace  J Am Ceram Soc. 2018;101:2638-2652.  \\x0c', 'ZrOxC(1\\x00x)  layer9,10,22 prior  to the  formation  of  an oxide  layer.  This  oxycarbide  layer  has,  however,  never  been  experimentally observed, perhaps because  the ZrC crystal  structure can accommodate up to 60% of oxygen crystal lattice without changing its structure.22 The overall reaction proposed by Katoh et al9 and Rama Rao and Venugopal22 is summarized as:  in  the  ZrC þ 1 2  ð1 \\x00 xÞO2 ! ZrCxO1\\x00x þ ð1 \\x00 xÞC  (1)  ZrCxO1\\x00x þ 1 2  ð1 þ 3xÞO2 ! ZrO2 þ ðxÞCO2  (2)  What is missing from the model et al9 are the effects of  proposed  by Katoh  formation of  the different  zirconia  polymorphs  and the  influence of  carbon oxidation on the  oxide microstructure with temperature. Carbon is produced during either ZrC oxidation or ZrOxC(1\\x00x) oxidation and is considered to remain in the oxide layer as inclusions at from 653 < T < 873, or to react with leaving the system as CO/CO2.9 characterization between the carbide  low temperatures,  oxygen, above 870 K,  The  interface  and  the oxide on samples oxidized at 1073 K and the Maltese  cross shape development was publication.23  the  topic  of  our  previous  The Maltese  cross  shape  of  the  oxide  is  induced by crack development at sample corners during the reaction.23 The  initial  stages of  the  failure of  corners pro motes  rapid growth of  the oxide in this  region. The inter face  between  the  carbide/oxide  of  a  hot-pressed  ZrC  specimen oxidized at 1073 K for 1 hour in air was comprised of a 2-lm-thick amorphous carbon matrix with ZrO2 nanocrystals embedded in it. Based on this finding, the pre sent work  gives  deeper  understanding  of  the  role  of  this  interface  in the reaction mechanism from 1073 to 1373 K  for hot-pressed ZrC specimens. Macroscopic  characteriza tion via XRD and SEM of  the oxide layer at different  tem peratures was coupled with the use of an isotopic exchange  technique where  oxygen  diffusion was (18O).  tracked  by SIMS  analysis using an isotopic tracer  2  |  E X P E R IM E N T A L P R O C E D U R E  High-density ZrC pellets were made using the same procedure described in ref,23 where the description of Set A, B, and C samples can be found.23 Commercial ZrC powder (3-5 lm; 90% < 8 lm with 0.2% < Hf < 2%; Grade B; H.C. Starck, Karlsruhe, Germany)  has,  according  to  the  manufacturer,  a  carbon  concentration  of  11.6 wt% (with  11.1% being combined and 0.5% being free carbon). Set B  ZrC pellets made from this powder in this study had a carbon concentration of 11.9% with an error ≤0.1% measured with an EMIA 320 V2 carbon and sulfur chemical analyzer  (Horiba  Scientific, Kyoto,  Japan).  The  commercial  ZrC  powder analyzed with the same unit showed a carbon concentration of 11.7% with an error ≤0.1%, B pellets acquired additional carbon on hot pressing. Set A  showing that Set  ZrC pellets presented a carbon content of 11.2%. The same  technique was used to quantify the carbon concentration in  the oxide layer.  Samples  from Set A, hot pressed at 2123 K, were used  for characterization of  the ZrC and ZrO2 layers across their interface with a Focused Ion Beam instrument coupled with  a Secondary  Ion Mass Spectrometer  (FEI FIB200-SIMS;  FEI Company, Hillsboro, OR).  This microscope  uses  a  gallium source for  imaging and sputtering. A high current  (20 000 pA) was used to sputter  the sample cross-sectional  surface for 5 minutes, whereas imaging was performed at a  current between 1000 and 3000 pA. SIMS chemical analysis was done on a 100 lm square region made of 10 rectangles of dimensions 10 lm 9 100 lm across sputtered at a current of 3000 pA. From each rectangle, sec the  interface  ondary negative and positive ions were collected to evaluate  the chemical abundance of carbon and oxygen species. On  each rectangular area previously analyzed by SIMS, energy dispersive X-ray spectrometer  (EDX) chemical analysis was  also  performed with  the  INCA system (INCA; Oxford  Instruments, UK) equipped with an ultrathin polymer win dow in  a LEO Gemini  1525  FEGSEM (Zeiss, UK),  at  20 kV. To evaluate the amount of material  removed by the  action  of  the  FIB-SIMS  a  nondestructive  technique was  used:  the Zygo NewView 200  3D optical  interferometer  which  enables  surface  topography  characterization.  The  sputter depth of the gallium source was approximately 0.6 lm in the ZrC side and 1.6 lm in the ZrO2 side. Set B, 1 cm3 cubes, hot pressed at 2273 K, were used  for kinetic analyses and oxide layer growth measurements.  Chemical  characterization with EDX on sample cross  sec tions polished and coated with chromium was performed in  a  JSM 6400 (JEOL, USA). Characterization of  the oxide  layer was performed by X-ray diffraction (XRD) on a Bruker D2 Desktop (MA, USA), using a Cu Ka source. XRD  was performed on oxide  fragments or on powders, which  were prepared by crushing the oxide  layer with an agate  pestle  in  an  agate mortar. Quench  treatments  for  kinetic  analyses  in  a  chamber-lift  furnace were  performed  in  air  atmosphere at 1073, 1173, 1273, and 1373 K at  times from  15 until 480 minutes. Prior  to insertion in the furnace,  the  samples were weighed and placed in a zirconia or alumina  crucible  so  that  the monitoring  of weight would  not  be  affected  by  oxide  layer  spallation. At  1073  and  1173 K  oxidation gave partly oxidized samples and the oxide layer  measurements were  performed  in  an  optical microscope  (Olympus UC30, Tokyo,  Japan) on sample  cross  sections  after  cutting  the  samples  in  half with  a  diamond-copper  saw (Isomet  Low Speed  Saw;  Buehler,  Braunschweig,  GASPARRINI ET AL.  |  2639  \\x0c', 'Germany). Oxide thickness was measured by averaging the  measurements  taken in the middle point of each side. The  oxide  layer  in the  samples oxidized at 1273 and 1373 K  did not adhere to the carbide core but detached during han dling,  therefore the oxide thickness was estimated from the  residual  thickness measured using electronic calipers (accu racy  of  0.01 mm)  in  the  centre  of  the  surface. Weight  change  and  surface  area measurements were  recorded  on  samples before and after quenching. Kinetic evaluation was  performed first by plotting the change in weight normalized  per  initial  surface  area and then by plotting the change  in  weight normalized per carbide core surface area of  the car bide after quenching as a function of  time. Data were then  fitted according to Equations 3 and 4:  DW  A  ¼ kl t \\x132¼ kp t  (3)  DW  A  \\x12  (4)  where DW is the weight change after oxidation (g), A is the surface area (m2), kl and kp are the linear and parabolic (g2 m rate constants, respectively, and the time (seconds).24,25 The mass gain normalized per  (g m  \\x002 s  \\x001)  \\x004 s  \\x001)  t  is  initial surface area takes into account  the initial area of  the  manufactured ZrC sample, whereas  the mass gain normal ized per carbide surface area takes into account  the area of  the carbide sample core evaluated at each quench time. For  experiments at 1273 and 1373 K the carbide core area was  measured with the calipers as  the oxide was not adherent.  The  area  in  this  case  is  evaluated with Equation 5  or  6  depending on the  shape of  the  carbide  core which turned  from a cuboid shape to a spherical one (see results in our previous publication23). Acarbide ¼ 2ða \\x01 bÞ þ 2ðb \\x01 cÞ þ 2ðc \\x01 aÞ  section  (5)  Acarbide ¼ 4pr 2  (6)  where a, b, and c are the dimensions of  the cuboid and r in  Equation 6 is the radius of  the sphere. The carbide area for  samples  oxidized  at  1073  and  1173 K was  evaluated  by  subtracting the dimensions of  the partly oxidized cube mea sured in the centre of each face using caliper from the oxide  layer thickness evaluated with the optical microscope. For long oxidation times (>2 hours) the oxide layer grew drasti cally assuming the typical Maltese cross shape so the oxide  thickness was measured in the center of the face of the cross  section as this layer showed a compact structure.  In  situ XRD analysis was  performed  on  commercial  ZrC powder using a high-temperature XRD (HT-XRD); X’pert Multi-Purpose Diffractometer terdam, the Netherlands) under air. Oxidation was  (MPD; Philips, Ams simu lated  by  heating  the  sample  in  air  on  a  platinum holder  (HDK 2.4; Buhler, Germany) at 40 kV and 40 mA, using Ni-filtered CuKa radiation between room temperature 1306 K, with scans performed every 50 K.  and  Set C, cut  from a Set B sample, was used for a double  oxidation  experiment  using  an  16O  followed  by  18O  enriched atmosphere  in an isotopic  exchange  set up. The  sample was  placed on a \\x007 mbar, inserted when a temperature  crucible  inside  a  quartz  furnace  evacuated  to  10  a  flux  of  200 mbar  of  oxygen  was  of  about  1060 K was  reached. The sample was first exposed to the oxygen atmosphere for 4 minutes in research-grade 16O (mostly consist16O plus 18O abundance of 0.2%),  ing of  the natural  then  the  furnace was  evacuated  and  cooled  before  the  second  oxidation was performed in an atmosphere of 40% enriched 18O gas at 200 mbar at  the same temperature and time as  the first oxidation.  FIB-SIMS was used to mill a ramp across the ZrC/ZrO2 interface using an Oregon-Physics Plasma FIB (Hillsboro,  OR) 500 nA/30 keV/Xe beam. This ramp had a 26.6° and width of 150 lm. The roughness of the top surface was 5.2 lm. SIMS and EDX analysis (done in a LEO  slope of  Gemini  1525 FEGSEM,  at  20 kV) were  performed  on  a  ramp made with the FEI FIB200-SIMS in the center of  this region. This final ramp used for analysis showed a 30° slope with a width of 45 lm and height of 40 lm. This ramp presented a smoother surface compared to the ramp  milled by the Plasma FIB so that  the SIMS data collected  were less noisy. Point scan EDX was performed and errors  were evaluated by averaging the values and taking the stan dard deviation of  the measurements done on the same area.  A MATLAB code was used to plot  the SIMS data on the  ramp. tool26  18O  \\x00  isotopic  fraction,  called  here  IF,  is  a  useful  as  it  represents  the  oxygen  isotope  distribution  on  the surface and limits any sharp change in ion signal intenlocalized charging.27  sities due to sample topography or  IF  is evaluated with Equation 7:  IF ¼  18O  \\x00  16O  \\x00 þ 18O  \\x00  (7)  where  16O  \\x00  and 18O  \\x00  represents  the oxygen ions  intensi ties detected by FIB SIMS.  3  |  R E S U L T S A N D D I S C U S S I O N  3.1  | Oxidation kinetics  Figure 1A shows  the mass gain normalized per  initial  sur face  area  vs  time  of  samples  quenched  from 1073  to  1373 K. Data initially were fitted to a parabolic law (Equa tion 4) as  shown in Figure 1B. This  first analysis  is, how ever, misleading  as  it  does  not  take  into  account  the  transformation  of  the  carbide  core  surface  area  evolution  which  is  known  to  be  of  primary  importance  in metal  2640  |  GASPARRINI ET AL.  \\x0c', 'oxidation.28 We  consider  this  to be  even more  important  for oxidation which involves a Maltese cross  evolution of  the oxide as it affects dramatically the carbide core surface  area. Figure 1C shows  the mass gain normalized per  car bide  core  area vs  time. Data  for oxidation at 1073 K can  be  fitted  to  one  kinetic  regime, whereas2  regimes  are  needed for  samples oxidized at higher  temperature  (1173 1373 K). The kinetics at 1073 K are linear, whereas  those  at 1173, 1273, and 1373 K fit a linear  trend up to 2 hours  of  oxidation  (shown  in  Figure 1D)  followed  by  a  rapid  change  as  the  oxidation  growth  accelerates  (Figure 1C).  Oxidation at an accelerated rate is referred to as breakaway  oxidation and is usually accompanied by loss of a proteclayer.29  tive  oxide  In  this  case,  breakaway  oxidation  occurred when the Maltese  cross  shape of  the oxide  fully  developed. The oxide growth in this region is severe and is  accompanied by shrinkage of  the unoxidized carbide  core  that changes  from its initial cubic shape to a spherical one et al23). For oxidation times over 2 hours  (see Gasparrini  the increase in oxidation was linked to the shrinking of  the  carbide  core which was  found in the form of a sphere, at  the  same  time  the  oxide  develops with  the  characteristic  Maltese cross shape.  The drastic change in the carbide core from a cubic into shape, described in our previous publication,23  a spherical  is explained with the fact  that corners are the points where  the oxidation proceeded more  rapidly developing with the  typical Maltese cross  shape. The carbide core surface area  F I G U R E 1  Plots of ZrC samples oxidized in 1 atm of oxygen at 1073-1373 K quenching at 900-1800-3600-7200-14 400-21 600-28 800 s.  A, weight change normalized per  initial surface area; B, square of weight change normalized per  initial surface area showing parabolic fitting; C,  weight change normalized per carbide core surface area showing 2 kinetic trends:  linear up to 2 h (7200 s)  followed by rapid oxidation for  samples oxidized from 1173 to 1373 K; D,  linear  fitting of  the data up to 2 h shown in C [Color  figure can be viewed at  wileyonlinelibrary.com]  GASPARRINI ET AL.  |  2641  \\x0c', '2642  |  is  considered  the most  suitable  parameter  to  be  used  to  derive kinetic analysis when compared with the initial  sur face  area,  compare Figures 1B vs  1D,  as  data  points  in  GASPARRINI ET AL.  Figure 1D showed very good agreement with a ting, R2 ≥ 0.99. From the in Figure 1D, it is possible  fitting of  evaluate  linear  to  linear  fit the data plotted  the  linear  rate  F I G U R E 2  A, Plot of  the linear  rate constants (kl)  from Figure 1D as a function of  temperature (T); B, plot of  the oxide layer  thickness vs  time for ZrC oxidation at 1073-1373 K in air; C, plot of  the oxide thickness linear  rate constants (kl,oxide) evaluated from the slope of straight  lines in B; D, XRD characterization of ZrC hot-pressed specimens oxidized at 1073-1373 K in air, scans were performed either on top or bottom the oxide layer: peaks are labeled as ■ for m-ZrO2 or ● for (left) oxidized at 1073 K for 4 h showing the points where the oxide layer was measured (yellow arrows),  t/c-ZrO2 or ZrO1.95.; E, photo of a partly oxidized ZrC sample cross section:  (right) oxidized at 1373 K for 30 min  of  showing the gap between the carbide and the oxide; F, sample shown in E (right) highlighting the points where XRD scans were done:  top side  (top ZrO2 surface) and bottom side (black layer once in contact with ZrC)  [Color  figure can be viewed at wileyonlinelibrary.com]  \\x0c', 'constants,  kl, which rate constants do not show a clear  are  shown  in Figure 2A. The  linear  trend: oxidation at 1073  and  1173 K showed  an  increase  in  oxidation  rate  but  at  1273 K a  slowdown  occurred  as  the  kl  is  lower  than  at  1173 K. This  slowdown in oxidation kinetics  is  accompa nied by an increase  in oxide  thickness  above 1273 K,  as  can  be  seen  in Figure 2B,C. Figure 2B shows  the  oxide  thickness vs  time  from 1073 to 1373 K,  the  slope of  the  straight  lines  fitting the oxide  thickness  rate  are  shown in  Figure 2C and  they  are  called  the  oxide  thickness  linear  rate constants kl,oxide. The oxide thicknesses have a similar growth rate at 1073 and 1173 K, however, at higher tem peratures the rate of oxide layer growth increased consider ably.  The  slowdown  in  kinetics  at  1273 K  which  corresponds  to  an  increase  in  oxide  layer  thickness was  investigated through experimental observations. Specimens  oxidized at 1073 and 1173 K presented an adherent oxide  layer, which withstood  cutting with  a  diamond  saw (see  photo  in  Figure 2E left).  Samples  oxidized  at  1273  and  1373 K instead  showed  a  voluminous  oxide  layer which  was easily damaged during handling. The oxide layer was  detached  from the  carbide  core  as  a  gap was  present  between the oxide and the carbide (see photos in Figure 2E  right and Figure 2F).  XRD analysis was performed on the bottom side (once  in contact with ZrC)  and top side  (exposed to air) of  the  oxide  layer  in samples oxidized at 1273 and 1373 K for  30 minutes as well as on the top surface of  the oxide layer  formed at 1073 and 1173 K for 30 minutes which was still  adherent  to  the  sample. XRD analysis  on  these  samples  showed that  the oxide  layer  formed at 1273 and 1373 K  was mostly made of m-ZrO2, whereas at 1073 and 1173 K the oxide was made of m-ZrO2 and either t/c-ZrO2 (see  Figure 2D).  A fine  black  powder was  found  between  the  carbide  and oxide  after oxidation at 1273 K and 1373 K and this  could either be  free  carbon arising from the  reaction hap pening at the interface, as previously reported by Voitovich and Pugach,16 or substoichiometric zirconia formed by anion vacancies30 reported to be black31 which presents the 2h = 30°  peak  at  observed  in  Figure 2D.  The  detached  oxide  layer  had  a  black  appearance  on  the  bottom side,  where it was once in contact with ZrC, and a white appear ance on the top side. This is in agreement with the observation of Shimada et al32 on ZrC single crystal oxidized at  from 0.02 to 2 kPa partial pressure of oxygen, however,  in  our  case  the black layer  could not be  separated from the  white  layer.  In  some  cases,  the  oxide was  cracked  and  lifted in the centre as  if  some gases pushed to make their  way out of  the bulk,  likely CO and/or CO2. The possibility that the oxide at 1073 and 1173 K could  be comprised of either  tor c-ZrO2 phase in the system is m-ZrO2 which  is due to the fact  that  the most  common  matches  all  the XRD peaks  present in tor c-ZrO2 and ZrO1.95 plus other peaks, apart from one at 2h = 30° which unfortunately is common with these phases (see PDF 037 148433 for m-ZrO2, 050 108934 t-ZrO2, 049 164235 and 01 081 154436 for ZrO1.95). An experimental oxidized at 1073 and 1173 K  for  for  c-ZrO2, observation  on  samples  showed  that  the  oxide  formed was  compact  and  stress  resistant. For  example, during a TGA/DTA experiment  at  1073 K,  a  partially  oxidized  sample  underwent  volume  expansion  and  cracked  the  alumina  crucible without  any  sign of crumbling. As a result it is plausible the peak at 30° is from tetragonal zirconia which is known to improve  the  strength  and  toughness  of m-ZrO2 when within it due to its ability to undergo a stress-induced polytransfomation.37 et al,22  dispersed  morphic  Rama  Rao  however,  detected  the  cubic  polymorph  of  zirconia  below 1073 K  during ZrC powder  oxidation which  then  transformed  to  the monoclinic polymorph at higher  temperatures. At equi librium, pure c-ZrO2 than those investigated here, 2950 K,38  is  stable at much higher  temperatures  between  2640 K up  to  the  melting  point  at  but  other  authors  found  the  cubic polymorph during oxidation of ZrC at temperatures.39 The identification of  relatively low  the second polymorph  in this work is considered to be either  tand/or c-ZrO2 due found in literature and the difficulties  to the discrepancies  of  differentiating  these  2  polymorphs with  the  available  techniques  (XRD, TEM) when the monoclinic polymorph  is also present. Still  it  is evident  that  the second polymorph  found in samples oxidized at 1073 and 1173 K enhances  the capabilities of zirconia to sustain stresses.  XRD analysis on the crushed oxide layer  formed during  oxidation at 1273 K and 1373 K showed this  to be  com prised of m-ZrO2 der), however, when XRD analysis was performed on the  (see Figure 2D labeled as 1273 K pow bottom and top sides of  the oxide  the XRD pattern chan ged. The bottom/black surface was  comprised of both m ZrO2 and either t/c-ZrO2 or substoichiometric ZrO1.95. The black color can be explained either by carbon traces or by  the presence of an oxygen-deficient m-ZrO2 layer, identified by Sinhamahapatra et al31 as ZrO2\\x00x. Substoichiometric zirconia could be difficult to imagine in a highly  enriched oxygen atmosphere such as atmospheric air.  It has  been, however, previously reported during zirconium oxi\\x008 Torr by found by Ma  dation in an atmosphere of oxygen of 2 9 10 Ma et al.40 et al,40  The  substoichiometric  layer  called ZrOx, was and stoichiometric  sandwiched  between  zirconium  metal  ZrO2. made of Zr/ZrOx/ZrO2 was considered energetically more stable to form than an interface Zr/ZrO2.40 to ZrC was identified as  This  3-layered  structure  In  our  study,  the black layer next  substoichio metric zirconia because it well matched the ZrO1.95 XRD pattern and its black appearance31 could only be assigned to the presence of ZrO2\\x00x and not  to carbon traces. Carbon  GASPARRINI ET AL.  |  2643  \\x0c', 'analysis was  indeed performed on both the top white side  and  the  bottom black  side  of  the  sample:  both  analyses  revealed  the  same  carbon  content. Carbon  analysis  per formed on a  fragment  taken from the white  top side of  a  sample oxidized at 1273 K for 8 hours 0.42 \\x06 0.01%, whereas sample oxidized at 1273 K for 0.42 \\x06 0.03% carbon. 8 hours was chosen so that it was possible to measure the  revealed a  carbon  content  of  the  oxide  black  layer  from a  30 minutes,  con tained  The  sample  oxidized  for  carbon content  (%) on the  top side of  the oxide  layer  far  from the ZrC side. The  sample  oxidized  for  30 minutes  was chosen to have enough black layer  to perform multiple  analyses. The  top side of  the oxide was  comprised of m ZrO2 as can be seen in Figure 2D. At higher 1273 K, where the oxidation kinetics is slowed, we see that  temperatures,  the intermediate layer  is missing and the oxide next  to this  layer  is  substoichiometric  zirconia. At  the  same  time,  the  oxide thickness rate increases (see Figure 2C). The volume  expansion from cubic to tetragonal and tetragonal +4%,42  to mono clinic  is  approximately  +2%41  and  respectively,  therefore  the  voluminous  and  porous  appearance  of  the  oxide which grew thicker at 1273 and 1373 K compared to  that  formed after oxidation at 1073 and 1173 K could be  explained by the stresses induced by these transformations.  The apparent  slowdown in the oxidation kinetics, kl (mea\\x001) observed at 1273 K (see Figure 2A) related to either a mass gain rate decrease or a  sured in g m  \\x002 s  could be  surface paper23  area  increase.  Results  shown  in  our  previous  showed that  the  surface  area decreased over  time  during  oxidation,  therefore  the  slowdown  in  kinetics  is  related to a mass gain rate decrease.  In carbides oxidation  there  is  a  competition between carbide oxidation and car bon oxidation (see Equations 8 and 9). ZrC þ O2 ! ZrO2 þ C  (8)  C þ O2 ! CO2  (9)  During uranium carbide (UC) oxidation,43 carbon oxidation occurs simultaneously with UC oxidation,44 however, in the case of ZrC, previous work18  suggests  that CO/CO2 In reference18  is produced mostly during the cooling stage.  the cracked sample surface was associated with the thermal  expansion  coefficient mismatch  between  the  carbide  and  oxide  polymorphs  during  the  cooling  stage.  However,  cracks were observed during the heating stage of the reaction using an in situ technique on ZrC23 and UC.43 These  cracks would offer a route for CO/CO2 to leave the sample during the heating stage. The mass gain rate decrease  observed at 1273 K could be related to the loss of carbon  due  to carbon oxidation (Equation 9) over  carbide oxida tion  (Equation 8).  This  occurs when  reaction 9  prevails  over  reaction 8 or when carbon already present  in the sys tem is oxidized following reaction 9. Carbon oxidation is  thought  to prevail in the intermediate T = 1273 K following  layer between ZrC/  ZrO2 samples oxidized at 1073 K (where  for  the  results  achieved  on  a  layer of  amorphous  carbon was detected at the interface carbide/oxide, see Gasparrini et al23). The gap observed between the carbide and  oxide  for oxidations  at 1273 and 1373 K could therefore  be  induced  by  preferential  oxidation  of  carbon. Oxygen,  when in contact with the intermediate layer made of carbon  and zirconia, preferentially oxidizes  the carbon that  leaves  the  sample  as CO or CO2 via as seen by Bellucci  cracks through the porous et al.18 The  oxide  layer,  preferential  oxidation of carbon promotes an oxygen-deficient zirconia study with PDF 01 081 154436  layer, characterized in this  for ZrO1.95, ichiometric zirconia is only found on the oxide next  to remain in the carbide/oxide gap. The substo to the  carbide in little amount, hence why it can only be detected  when XRD is directly performed on the black surface of  this oxide layer.  3.2  |  In situ HT-XRD  Formation of t-/c-ZrO2 below the generally accepted polymorphic stability ranges which start at 1448 K45 for t-ZrO2 and 2643 K45 confirmed during an oxida for  c-ZrO2 was tion experiment performed on ZrC powder  in a high-tem perature XRD. When ZrC powder  is heated in a static air  environment  from room temperature  to  1306 K (see Fig ure 3)  the  first oxide polymorph formed matched the  t-/c ZrO2 t-/c-ZrO2 XRD pattern was 891 to 1203 K, whereas for temperatures above this  pattern. The  seen  from  range  m-ZrO2 peaks appeared (XRD peaks were matched with PDF 037 148433 for m-ZrO2, 050 108934 for t-ZrO2, 049 164235 for c-ZrO2 and PDF 035 078446 for ZrC). There are several ways in which t-ZrO2 can be including critical crystallite size and the  stabi lized over m-ZrO2 presence of ion impurities or oxygen vacancies. Regarding  the  critical  crystallite  size,  the  formation of  t-ZrO2 before m-ZrO2 can be explained with the theory of Garvie47,48 which states that a “critical crystallite size” exists for ZrO2. Below a size of 5-16 nm t-ZrO2 crystals are favored over m-ZrO2 crystals, whereas above a critical size of 30 nm mZrO2 crystals are favored. This occurs because t-ZrO2 crystals are stabilized due to their reduced crystal surface free energy49 when showed23 that  compared  with m-ZrO2. the crystallites of ZrO2 formed were approximately 10 nm in size which agrees well with the theory of Garvie.48  Previously we  Another  mechanism  proposed  by  different  authors  for  the  stabilization  of  t-ZrO2 tures, even down to room temperature, is oxygen ion vacancies.45 The presence of  at  lower  tempera the presence of  stable  tetragonal  phase  in  bulk ZrO2 vacancies within the ZrO2 lattice supposedly generated temperature or under vacuum.45 Srinivasan  has  been  explained  by  oxygen  ion  either at higher  2644  |  GASPARRINI ET AL.  \\x0c', 'et al50  reported that  the  transformation from tetragonal  to  monoclinic ZrO2 oxygen at relatively low temperature,  is  related  to  the  adsorption  of  gaseous  such as 573 K or on  cooling during an HT-XRD study of zirconia.  In our case  the transformation from tetragonal  to monoclinic occurs at  high  temperature,  above  1099 K shown  in  Figure 3  and  could  be  related  to  crystal  growth or  removal  of oxygen  vacancies  from the ZrO2 substoichiometric  lattice. The  characterization of  a  layer of  zirconia on hot-pressed ZrC in  samples oxidized at 1273 and 1373 K shown in Figure 2D  suggests that  the first oxide formed during ZrC oxidation is  characterized by oxygen vacancies.  Previous studies of ZrC oxidation,  typically identify the  ZrO2 which, instead, is typically identified in amorphous ZrO2 oxidation studies.47,52 Due to this controversy and the simi crystallites  as  cubic22,39,51  rather  than  tetragonal  larities in their XRD patterns, as previously explained, both  tetragonal  and  cubic  polymorphs may  be  present  in  our  system.  3.3 Chemical analysis across the ZrC/ZrO2 interface  |  A sample  cross  section (shown in Figure 4A)  from set A  oxidized in a  furnace  for 1 hour  at 1073 K was  analyzed  by  FIB-SIMS  and  SEM.  Chemical  analyses were  per formed by SIMS and EDX on each rectangle with size 100 9 10 lm highlighted in Figure 4A. SIMS is a surface characterization technique, whereas EDX gives information  on the sample bulk. Figures 4B,C, and D show the normal ized ions counted during SIMS analysis which are relevant  to this system (expressed in fractions of counts per second:  cps/cps):  (B)  carbon,  (C)  oxygen,  and  (D)  zirconium related species. SIMS analysis showed that on the ZrC side  no oxygen was present  and in the ZrO2 (see Figure 4B,C). In the same way case ZrC2(+), (see Figure 4D) and dramatically side just 10 lm away from  side  little  carbon  was  found  the  ions  related  to ZrC,  in  this  showed  the  highest  yield  in  the ZrC side  decreased to zero on the ZrO2 the carbide/oxide interface.  EDX performed on the same region after  the FIB-SIMS  analysis is shown in Figure 5A. EDX was done by selecting each 100 9 10 lm rectangle as previously done with SIMS, however, it was also performed outside the FIB SIMS sput tered region. Chemical analysis was performed considering  all elements: no oxygen was found in the ZrC area and no  carbon was  found in the ZrO2 sputtered region as shown in Figure 5A (EDX peaks within  side within the FIB-SIMS  this  region are shown in Figure 5C). EDX performed out side the sputtered region reveals oxygen in the ZrC region  and carbon in the ZrO2 was used for FIB sputtering, however, it is neglected in the analysis as its abundance was ≤0.54 atomic (%).  region. Gallium was present  as  it  The presence of oxygen in ZrC and carbon in the ZrO2 side outside the FIB sputtered region could be related to  contamination due either  to the presence of an oxide layer  (the  sample was  kept  in  air  atmosphere)  or  to  sample  preparation effects  (such as  the  effect of polishing or  the  chromium coating). Both SIMS and EDX analysis,  how ever,  are  in  agreement within  the  FIB-SIMS  sputtered  region that  showed the presence of only ZrC and ZrO2 on the left and right side of the interface region. EDX is well known to lack sensitivity for  light elements, such as carbon  and  oxygen,  however,  the  similarities  between  the EDX  and  the SIMS analyses  give EDX results  credibility  and  this is why EDX is included in this work.  BSEI of  sample cross  sections  (Figure 5B)  show the 3  regions: ZrC (zone  1),  intermediate  layer  (zone  2),  and  ZrO2 1073 K were analyzed in BSE mode,  (zone  3).  Two  samples  from Set  B  oxidized  at  the intermediate layer  on the sample oxidized for 30 minutes was approximately 0.7-6.5 lm wide, whereas sample oxidized for 4 hours at the same temperature was approximately 1.3-9.8 lm wide. An increase mediate layer width occurs with time, however,  the  intermediate  layer  in  the  in the  inter this has an  irregular  thickness  (see  Figure 5B)  and  so  the width  is  expressed as a range. EDX analysis  for a sample oxidized  for 30 minutes  at 1073 K is  shown in Figure 5D; oxygen  was deliberately omitted in the  analysis done on the ZrC  side (zone 1) as the oxygen was found to be related to sur face  contamination  in  the  analysis  shown  in  Figure 5A.  Chromium was also neglected in the analysis as  the abun dance of 0.51% arises  from the  chromium coating depos ited  during  sample  preparation.  Two  analyses  were  performed and the results are summarized in Table 1.  F I G U R E 3  HT-XRD of ZrC powder oxidized from room  temperature to 1306 K on a platinum foil with scans taken every  50 K under static air. The last scan was taken on the powder cooled to room temperature. Peaks were indexed using PDF 035 078446 for ZrC, PDF 037 148433 for m-ZrO2, 050 108934 for t-ZrO2, 049 164235  for c-ZrO2 or compared to the platinum holder blank XRD pattern  [Color  figure can be viewed at wileyonlinelibrary.com]  GASPARRINI ET AL.  |  2645  \\x0c', 'EDX analysis shows that Zr and O are present  in a stoi chiometric ratio 1:2 in the oxide layer but also in the inter mediate  layer  region. This  result  is  in agreement with the  observations  performed  via  transmission  electron micro scopy where the  intermediate layer was  found to be com prised of ZrO2 nanocrystals embedded carbon layer.23 EDX also shows that  in  an  amorphous  the intermediate layer  is comprised of carbon and ZrO2.  3.4  | Oxygen tracer diffusion  An  oxygen  tracer  technique was  used  to  determine  how  oxygen  combines with  the  carbide  during ZrC oxidation.  Figures 6A,C show the top view of  the sample surface and  the attempts made to create a ramp across  the oxide layer  suitable for SIMS analysis. The only successful  ramp used  for  this  study is highlighted and shown from the  side  in  Figure 6B.  Figures 6D,E show a BSEI and a high magnification ion  image of  the ramp across  the ZrC/ZrO2 and polished by FIB. Figure 6E is the region used for SIMS  interface sputtered  analysis. The 3 regions observed: ZrC,  intermediate  layer  and ZrO2 chemical  are  labeled  as  1,  2,  and  3,  respectively. EDX  characterization  on  these  regions  is  shown  in  Table 2. Figures 6F,G show the SIMS elemental mapping:  Figure 6F shows  the  total  oxygen  abundance map  across  this  region, whereas Figure 6G shows  the  18O  \\x00  isotopic  fraction (IF) distribution on the same area. Figure 6F shows sum of both 16O and 18O  that  the total oxygen,  \\x00  \\x00  ions,  is  evenly distributed across zone 3 apart  from sharp changes  F I G U R E 4  A, SEI of a sample cross section oxidized at 1073 K for 1 h (set A), the squared region divided into 10 rectangles of size 100 9 10 lm is the region where FIB-SIMS analysis was performed followed by EDX analysis,  the dashed rectangle highlights the interface;  normalized values for  the (B) carbon-related secondary ions; C, oxygen-related secondary ions; D, zirconium-related secondary ions sputtered  from the surface and expressed in cps/cps (counts per second)  [Color  figure can be viewed at wileyonlinelibrary.com]  2646  |  GASPARRINI ET AL.  \\x0c', 'GASPARRINI ET AL.  |  2647  F I G U R E 5  A, EDX analysis performed on each rectangle shown in Figure 4A after FIB-SIMS analysis showing composition in atomic (%),  stripes show the stoichiometric values of ZrO2 and ZrC; B, BSEI of a sample cross-section oxidized for 30 min at 1073 K (set B),  red dashed  lines highlight  the interface region (dark zones)  labeled as 2, zone 1 is ZrC, and zone 3 is ZrO2, starred symbol highlights points where EDX  was performed; C, example of EDX spectra acquired from the ZrC and ZrO2 sides within the sputtered region; D, EDX analysis performed on  each point  labeled in B showing composition in atomic (%),  the different size of  the symbols refer  to the 2 EDX analyses with 2 symbols sizes  shown in B [Color  figure can be viewed at wileyonlinelibrary.com]  T A B L E 1  Chemical composition of  the ZrC side,  intermediate layer, and ZrO2 side (labeled as 1, 2, and 3 in Figure 5B) of sample oxidized  at 1073 K for 30 min. All data shown are in atomic (%)  ZrC side (1)  Zr (%)  37.6 \\x06 5.7 C/Zr = 1.6 \\x06 0.3  C (%)  61.5 \\x06 6.1  Intermediate layer (2)  Zr (%)  22.7 \\x06 0.7 O/Zr = 2 \\x06 0.2  O (%)  44.2 \\x06 2.4  C (%)  33.2 \\x06 3.0  ZrO2 (3)  Zr (%)  33.5 \\x06 2.9 O/Zr = 1.9 \\x06 0.2  O (%)  62.8 \\x06 5.2  C (%)  18.4 \\x06 1.7  intensities due to sample’s uneven morphology at cracks of and pores. Figure 6G reveals a gradient of 18O throughout is 5.8-16.9 lm wide  intermediate  (mea \\x00  zone 2,  this  layer  shown in Figure 6H. This method consisted in evaluating  an averaged  value  of  the  IF within  a  region,  highlighted  with a black square in Figure 6G, which was chosen as the  sured  from Figure 6E).  To  characterize  the  IF  gradient  border between zones 2 and 3 runs parallel  to the border  through the intermediate layer, an IF profile was computed  between zones 1 and 2. From the overall  IF profile across  by an integrated linescan method in MATLAB and this  is  zones  1,  2,  and  3  shown  in  Figure 6H,  the  IF  gradient  \\x0c', 'across  zone 2 was  isolated and highlighted with a box in  Figure 6H. This was normalized to 1 and is shown in Figfitted using the Fick’s of diffusion for a semi-infinite media, considering a constant source of gas27 (Equation 10):  ure 6I. This profile was  second law  C  0 ðxÞ ¼ CðxÞ \\x00 Cbg  Cg \\x00 Cbg  ¼ erfc  x  2  ffiffiffiffiffi  Dt  p  \\x12  \\x13  (10)  where C0 (x) the natural  is the isotopic fraction which is normalized to 18O isotopic  abundance, Cbg concentration Cg during  is  equal  to 0.2%  and  the  gas  tracer  the  second  anneal  is  equal  to  40%. The  parameter  x  represents  the  thickness  of  the  intermediate  region where  the  diffusion  profile is measured, this is highlighted with the black box in Figure 6H, where x = 0 represents the border between zone  3 and 2 within the rectangle shown in Figure 6G;  t  repre sents  the  time of  sample  exposure  to the oxidizing gases  and D is the diffusion coefficient. The obtained diffusion \\x001 with R2 = 0.93. coefficient D was = 9 9 10 for oxygen through this inter \\x0010 cm2 s  The diffusion coefficient  mediate  region  in  between  ZrC/ZrO2 coefficients  is  larger  than  the  grain-boundary diffusion mately DB = 1 9 10  which  is  approxi \\x0011 cm² s  \\x001  evaluated by Brossmann  (A)  (C)  (F)  (G)  (H)  (I)  (D)  (E)  (B)  F I G U R E 6  A, Top view image of ZrC sample oxidized in 16O/18O  the  environment: only the Plasma FIB ramp, 150 lm wide, was suitable for chemical  analysis; B, side view representation of  the  ramp; C, SEI of  the ramps milled on the  top surface of  the sample; D, BSEI of  the  Plasma FIB ramp, cracks are visible on the  oxide surface; E, high magnification ions  FIB image showing zones 1, 2, and 3 (ZrC,  intermediate layer, ZrO2,  respectively); F,  total oxygen map across the area seen in E; IF, 18O/(16O + 18O)  G,  isotopic fraction,  map,  the black dashed rectangle highlights  the area where the averaged IF was  calculated; H, averaged IF profile from the  region highlighted in G; I, IF profile from zone 2 fitted with the Fick’s law equation  [Color  figure can be viewed at  wileyonlinelibrary.com]  2648  |  GASPARRINI ET AL.  \\x0c', 'et al53 with undoped ultrafine monoclinic ZrO2 at 1073 K. The discrepancy with the value of Brossmann et al53 can  be  explained by the  fact  that  their ZrO2 than those present  crystals were  an  order of magnitude larger 70-300 nm53 compared to 10 nm (see TEM characterization in our previous paper23). The 18O diffusion coefficient  in our sample,  evaluated  in  this work,  however,  is  representative  of  the  oxygen diffusion through an intermediate layer where oxy gen not only diffuses but also reacts nanocrystals23 and/or CO/CO2. No 18O gradient is present  to either  form ZrO2  in zone 3 and therefore it  is  plausible that during the second oxidation step, either  18O  passed  through  an  almost  stoichiometric  zirconia  layer  without  significantly reacting with it or no diffusion gradi ent  is present here due  to the porous nature of  the oxide  layer. Confirmation of  the stoichiometric nature of ZrO2 EDX analysis on the oxide layer just a few  is  given  by  micrometers  away  from  the  intermediate  layer  (see  Table 2). The porous and cracked nature of  this oxide indiboundaries,23  cates  that  oxidation  proceeded  along  grain  and  cracks  are  clearly  visible  in Figures 6D,E. Closer  to  the  intermediate layer, zone 2, cracks  are  less pronounced  compared with the cracks seen by SEM on polished surfaces which extend to 20 lm from the carbide/oxide inter face  (compare  cracks  in  Figures 6E,D).  SIMS  analysis  suggests that oxidation proceeded with inward formation of  the oxide layer by oxygen penetration through cracks. The  cracks present  in the oxide layer mostly followed the grain et al23  boundaries  (see  Gasparrini  and  cracks  in  Fig ure 6D),  they are  induced either by tension related to the  volume expansion occurring in the transformation between  ZrC to ZrO2 or by stresses oxidizes into CO/CO2. The fact tion of ZrC are linear, (see Section 3.1),  related to gas effusion when C  that  the kinetics of oxida is  in agreement  with the observation of  the  intermediate  layer where oxy gen  diffusion  needs  to  take  place  before  this  carbon enriched layer can transform into ZrO2. The cyclic debonding of this intermediate layer evidenced in a previous paper23  shows  the paramount  role of  this  layer which acts  as a diffusion barrier  in the oxidation of ZrC at 1073 K in  oxygen. A schematic of what  is deduced to happen during  hot-pressed ZrC oxidation is shown in Figure 7.  At  1073  and  1173 K an  oxide made  of m-ZrO2 to the ZrC core  and  t and/or  c-ZrO2 developed adherent  via  an  intermediate  layer made  of  amorphous  carbon  and  nanocrystals of  zirconia  (this  is  schematically shown as  a  gray  layer  in  Figure 7  top  row between  ZrC  and  the  oxide;  the  characterization of et al23  this  intermediate  layer  is  shown  in  Gasparrini  in  a  sample  oxidized  at  1073 K). The oxide  layer developed via the  inward diffu sion of oxygen through this  intermediate  layer. Diffusion \\x001 using experiment in  of oxygen was measured to be 9 9 10  \\x0010 cm2 s  18O  as a tracer 16O/18O with  in  a  double  oxidation  SIMS. Diffusion  took  place  in  a  layer  of  constant thickness, as the oxide cyclically debonds after reaching a thickness of approximately 20 lm (see Gasparrini et al23), in agreement with the linear kinetic observed  at 1073 K (see Figure 1D). Diffusion is usually described  with a parabolic behavior,  in this case, however,  the over all  reaction rate was observed to be linear. Linear  reaction  rates can be ascribed to either  a surface reaction or a dif fusion mechanism through thickness.29 Oxidation  a  protective  layer  of  constant  at  1173,  1273,  and  1373 K fol lowed linear kinetic up to approximately 2 hours where  a  drastic  oxidation was  related  to  the ZrC core  change  in  shape from a cube to a sphere.  Increasing the temperature  to 1273 and 1373 K affected both the oxide morphology  and  the  intermediate  layer,  as  shown  in Figure 7  bottom  row. The  oxide  at  1273  and  1373 K was made  of m ZrO2 formed at 1073 and 1173 K as  and  it  was  voluminous  compared  to  the  oxide  the oxide  thickness  linear  rate constants  increased considerably (see Figure 2C). The  voluminous nature of  the oxide  at 1273 and 1373 K can  be  explained  with  the  volume  expansion  occurring  between t/c-ZrO2 and m-ZrO2 which is usually reported approximately 4%.42 Between ZrC and the oxide,  to be  a  gap  replaced  the  intermediate  layer which was  character ized at 1073 K. The oxide  surface  at 1273 K at  the gap  was  characterized as substoichiometric a yellow/gray ZrO2\\x00x between this region and  zirconia  (shown in  Figure 7 as  layer), with  no carbon  uptake  the  top  oxide  surface. \\x001), loss  The  decrease  in  linear  rate  constant,  kl  (g m  \\x002 s  at  1273 K compared  to  1173 K was  related  to  the  of  carbon. Carbon could leave the system as CO and/or CO2 due to preferential oxidation of carbon at the intermediate  layer,  reaction (9)  prevailing  over  reaction (8),  leaving  a  substoichiometric  layer of ZrO2 near  the  interface ZrC/  ZrO2.  T A B L E 2  Chemical composition of the ramp shown in Figure 6E: ZrC (zone 1), intermediate layer (zone 2), and ZrO2 (zone 3) for ZrC sample oxidized at 1073 K in an isotopic exchange setup for 4 + 4 min first in 16O then in a 18O atmosphere. All data shown are in atomic (%)  ZrC side (1)  Intermediate layer (2)  ZrO2 (3)  Zr (%)  O (%)  C (%)  Zr (%)  O (%)  C (%)  Zr (%)  O (%)  C (%)  45.1 \\x06 4.0 7 \\x06 2.5 C/Zr = 1.0 \\x06 0.1  46.2 \\x06 4.8  27.2 \\x06 3.2 42.8 \\x06 1.8 O/Zr = 1.6 \\x06 0.1  28.9 \\x06 3.0  34.0 \\x06 2.5 61.5 \\x06 4.5 O/Zr = 1.8 \\x06 0.2  15.8 \\x06 2.6  GASPARRINI ET AL.  |  2649  \\x0c', '4  |  C O N C L U S I O N S  The  key  role  of  the  intermediate  layer  between ZrC and  ZrO2 during ZrC oxidation has been investigated.  1. Oxygen diffusion at 1073 K through the compact  inter mediate  layer made  of ZrO2 phous carbon was revealed by a double oxidation experiment performed in 16O and 18O atmosphere on a  nanocrystals  and  amor ZrC hot-pressed sample. 2. The diffusion coefficient measured with the use of  18O \\x0010  as a tracer across the intermediate layer was 9 9 10 \\x001, R2 = 0.93. The cm2 s intermediate layer made  diffusion  of  oxygen  through  the  of  nanocrystals  zirconia  and  carbon  is  considered  the  rate-limiting  step  of  the  reaction.  3. This  intermediate  layer  was  observed  to  cyclically  debond when a 20 lm).  critical  thickness  is  reached  (approxi mately  The  overall  linear  kinetics  of  hot pressed ZrC samples  agrees with the oxidation mecha nism governed by oxygen diffusion through a  layer of  constant 4. Linear  thickness.  rate  constants  (g m  \\x002 s  \\x001)  showed a  slowdown  at 1273 K compared to 1173 K,  this was experimentally  correlated with the  loss of  the  intermediate  layer made  of  carbon.  Instead  of  the  intermediate  layer,  a  gap  between the  carbide  core  and the oxide was observed  on samples oxidized at 1273 K as well  as  at 1373 K.  Carbon  loss  at  the  intermediate  layer was  correlated  with carbon oxidation prevailing over zirconium oxida tion (Equation 9 prevails over Equation 8) which would  leave the intermediate layer as CO and/or CO2.  F I G U R E 7  Oxidation stages of hot-pressed (H.P.) specimens of ZrC: at 1073 and 1173 K the oxide made of m-ZrO2 and t/c-ZrO2 develops  attached to an intermediate layer comprised of amorphous carbon and ZrO2 nanocrystals. The oxide is adherent and well sustains stresses. At  1273 and 1373 K the kinetics is slowed down by carbon oxidation in the intermediate layer  that  leaves a gap between the carbide core and the  substoichiometric zirconia layer. The oxide layer  is made of m-ZrO2 and is voluminous, no adherent  intermediate layer  is observed [Color  figure  can be viewed at wileyonlinelibrary.com]  2650  |  GASPARRINI ET AL.  \\x0c', '5. At  the  gap  on  samples  oxidized  at  1273  and  1373 K  the oxide layer was black and it was  identified as  sub stoichiometric  zirconia  as  no  carbon  uptake  was  detected in the black oxide  formed at 1273 K next  to  the ZrC side  compared to the  top white  surface of  the  oxide layer at  the same temperature.  6.  Samples oxidized at 1273 and 1373 K showed a gap between ZrC and ZrO2\\x00x and the oxide was voluminous and made of m-ZrO2. On the contrary, the oxide formed by oxidation at 1073 and 1173 K strongly adhered to  the carbide core and was comprised of mand t/c-ZrO2. 7. The voluminous and heavily cracked structure of the  oxide at 1273 K and 1373 K demonstrates its nonpassi vating nature and explains  the oxide thickness  increase  at  these temperatures. Linear kinetics were followed by  a drastic  increase  in oxidation after 2 hours of oxygen  exposure for T = 1173-1373 K. This change was exper imentally  linked  to  the  carbide  core  having  its  initial  cubic shape evolving to a spherical one. 8. Cracking of the oxide and cyclic cracking of the interme diate layer is key for understanding the oxidation mecha nism of ZrC. An analytical model and a finite element  modeling (FEM)  approach will be developed to assess  the extent of stresses in the oxide during oxidation. 9. Rapid  oxidation  produced  voluminous  and  heavily  cracked  oxides  for  oxidation  from 1073  to  1373 K,  hence, ZrC is  not  suitable  for  applications  in  oxygen  environments in this range of  temperatures.  A C K N OW L E D GM E N T S  The  authors  are  grateful  to  the  EPSRC DISTINCTIVE  (Decommissioning,  Immobilization  and Storage  soluTIons  for  NuClear  wasTe  InVEntories)  Consortium for  their  financial  support  of  this  project  (EPSRC Industrial Case  Award EP/M507428/1 Grant  and  the DISTINCTIVE EP/  L014041/1 Grant). We  also thank Dr Mahmoud Ardakani  and Dr Ecaterina Ware for help with SEM-EDX analyses,  Richard Sweeney for help with HT-XRD, Dr Michele Pettin\\x12a for developing a finite element modeling code simulat ing oxidation of ZrC, Dr Catriona McGilvery for help in  characterize  t and/or  c-ZrO2 nanocrystals kalls for help with sample preparation.  and Garry Sta O R C I D  Claudia Gasparrini  http://orcid.org/0000-0001-8343-4733  Denis Horlait  http://orcid.org/0000-0002-2645-6896  R E F E R E N C E S  1. Paul A,  Jayaseelan DD, Venugopal S,  et al. UHTC composites  for hypersonic applications. Bull Am Ceram Soc. 2012;91:22-29.  2. 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},{
  "_id": 289,
  "PDF": "Zirconium Carbide Oxidation- Maltese Cross Formation and Interface Characterization.pdf",
  "Text": "['Oxid Met (2017) 88:509-519  DOI 10.1007/s11085-016-9672-6  O R I G I N A L P A P E R  Zirconium Carbide Oxidation: Maltese Cross  Formation and Interface Characterization  Claudia Gasparrini1 • Renaud Podor2 Denis Horlait1,3 • Richard Chater4 • William Edward Lee1  •  Received: 11 November 2016 / Published online: 2 December 2016 Ó The Author(s) 2016. This article is published with open access at Springerlink.com  Abstract Oxidation of dense hot-pressed ZrC specimens from 1073 to 1473 K was  investigated using an in situ technique: HT-ESEM. Cuboid specimens were moni tored on the surface and on edges and corners during oxidation in order  to under stand the influence of crack formation and propagation on the Maltese cross shape  development of  the oxide. The oxidation mechanism comprised three  steps:  (1)  delamination of sample edges,  (2) crack formation at corners and (3) crack prop agation towards the inner core and formation of microcracks parallel to the interface  that  increase the accessible surface area followed by a drastic volume expansion.  The microcrack pattern is  found to be repetitive as  if a cyclic debonding of  the  interface occurred. Characterization of the interface by TEM and HRTEM revealed  & Claudia Gasparrini  c.gasparrini14@imperial.ac.uk  Renaud Podor  renaud.podor@cea.fr  Denis Horlait  horlait@cenbg.in2p3.fr  Richard Chater  r.chater@imperial.ac.uk  William Edward Lee  w.e.lee@imperial.ac.uk  1  2  3  4  Centre for Nuclear Engineering (CNE) & Department of Materials, Imperial College London,  South Kensington Campus, London SW7 2AZ, UK  Institut de Chimie Se´ parative de Marcoule, UMR 5257 CEA/CNRS/UM/ENSCM, Site de  Marcoule, BP17171, 30207 Bagnols-Sur-Ce` ze, France  CNRS, Centre D’Etudes Nucle´ aires de Bordeaux-Gradignan, UMR 5797, Chemin Du Solarium,  33175 Gradignan, France  Department of Materials,  Imperial College London, South Kensington Campus,  London SW7 2AZ, UK  123  \\x0c', '510  Oxid Met (2017) 88:509-519  the interface between ZrC and ZrO2 to comprise a 2 lm thick amorphous carbon  matrix with ZrO2 nanocrystals embedded in it.  Keywords Zirconium carbide \\x01 Oxidation \\x01 Maltese cross \\x01 HT-ESEM  Introduction  Zirconium carbide (ZrC)  is a candidate ceramic for use in ultra-high temperature  applications  in hypersonic vehicles  [1, 2] and in the nuclear  industry as an inert  matrix fuel or as a structural component in tristructural-isotropic (TRISO) [3, 4] fuel  particles  for Generation  IV reactors.  Its  advantageous  properties  include  high  temperature stability in non-oxidizing atmospheres, high melting point (around 3700 K) [5], and high thermal conductivity (20 Wm-1K-1 at 300 K) [6]. However,  zirconium carbide  readily oxidizes  from relatively low temperatures  (as  low as  653 K according to Shimada and Ishii) [7]. A comprehensive understanding of the  oxidation mechanism and  kinetics  of ZrC is  thus  important  for  predicting  its  behaviour under accident scenarios in which ZrC would come in contact with air at  high temperature. Data reported in previous studies consider ZrC oxidation in the  form of powder [7-9], sintered pellet [10-13] and single crystal [14, 15] under many  different conditions. What is understood and well reported regarding ZrC oxidation is that ZrC is susceptible to rapid oxidation at high temperatures (T Z 870 K) [16]  and that  the mechanism of formation of the oxide occurs with the typical ‘‘Maltese  cross’’ shape [17], previously reported for other carbides and borides of group IV,  V, VI of the transitions metals [18-21]. However, an overall model describing the  controlling mechanism is difficult  to achieve as ZrC oxidation is strongly affected  by a range of parameters and previous  studies show limited experimental data in  overlapping  parameter  ranges  [16]. The  parameters  that mainly  influence  the  oxidation behaviour of ZrC are temperature, oxygen partial pressure, stoichiometry  (ZrC1-x) and samples  intrinsic properties derived from their preparation method  (i.e. porosity,  impurity nature and level  [16, 22]). An overall understanding of  the  ZrC oxidation mechanism is also difficult  to achieve as discrepancies in activation  energy values in experiments performed over the same range of temperature show  that  a  competition  between  diffusion-controlled  and  phase  boundary  controlled  reaction occurs [23, 24]. Kuriakose and Margrave studied the oxidation of electron  beam melted ZrC at 100 kPa of O2. For 827-925 K their measured kinetics were  found to follow linear rates throughout the temperature range studied while at higher  temperatures  (973-1123 K) a destructive oxidation was  reported. Shimada et al.  [25].  investigated the microstructure of the oxide after oxidation experiments from  773 to 873 K on ZrC single crystals. Observations of cross sections revealed two  oxide regions. Zone 1 was described as a pore free and compact intermediate region  enriched in carbon with crystallites of cubic zirconium oxide (c-ZrO2) that grew parabolically to 2-3 lm thick. This acted as a diffusion barrier against oxidation.  Zone 2 was instead described as a carbon-poor  layer of cracked zirconium oxide  which grew linearly above Zone 1. A mechanistic model of  the oxidation of ZrC  was first proposed by Rama Rao and Venugopal [23] and then by Katoh et al. [16].  123  \\x0c', 'The  proposed models  suggest  as  initial  stage  the  formation  of  an  oxycarbide  ZrOxC(1-x)  layer at  the interface [7, 16, 23]. This intermediate layer however has  never  been  experimentally  observed. The  present work  thus  aims  to  provide  information  and  understanding  on  the  oxidation mechanism of  hot  pressed  specimens of ZrC from 1073 to 1473 K with a particular attention to the nature  of  the intermediate  layer  and to the mechanism of  formation of  the oxide. The  peculiarity of  this work is  the use of  state-of-the-art characterisation techniques:  in situ high-temperature environmental scanning electron microscopy (HT-ESEM)  for  investigations on the Maltese cross  formation and TEM and HRTEM for  the  characterization of  the ZrC/ZrO2 interface.  Experimental Procedures  High density ([96% of  produced by hot pressing ZrC commercial powder  the theoretical density, TD) disc-shaped ZrC pellets were (3-5 lm, 90%  \\x0c', 'squared. General procedure and details about  the HT-ESEM platform can be found  in refs [26-28].  Microstructural characterisation of the carbide-oxide interface on partly-oxidised  cross sections was performed at  Imperial College London. First observations were  made with a SEM (LEO Gemini 1525 FEG-SEM, Zeiss,  Jena, Germany) using  secondary  electron  imaging  (SEI)  equipped with  an  energy  dispersive  x-ray  spectrometer  (EDX)  for  chemical  analysis  (INCA, Oxford Instruments, Oxford,  UK). Nanoscale analysis of the intermediate layer between the carbide and the oxide  was performed with a transmission electron microscope (TEM)  (2000FX,  JEOL,  Tokyo, Japan) and a High Resolution—TEM (JEOL JEM 2100F). Samples for TEM  analysis were prepared with the FIB (focussed ion-beam) technique (Helios Nanolab  600, FEI Company). Selected area diffraction (SAD) patterns were  indexed by  matching the dhkl values with reference patterns. A SingleCrystal software (version  2.3 CrystalMaker Software Limited, Oxfordshire, UK) simulation was performed to  recheck the assigned indexes. These characterisation techniques were performed on  samples cooled to room temperature. The results therefore show the structure of the  oxide after  it had undergone several modifications due to volume changes.  Indeed  ZrO2 undergoes transitions between monoclinic,  tetragonal and cubic polymorphs  within the range of temperatures used in this study [29, 30]. In particular the volume  change occurring in the transition from the tetragonal  to the monoclinic structure  during cooling is estimated to be about ?3 vol.% [31, 32]. To investigate the role of  crack formation in the oxide development as a Maltese cross excluding any effect of  cracking due to volume change when cooling,  the in situ experiment on a central  area of  the sample surface and on sample corner’s in the HT-ESEM is required.  Results and Discussion  Maltese Cross Formation  ESEM analysis  (Fig. 1) of  sample surfaces  shows  the transition from a compact  oxide  to  a  cracked  oxide  layer. Cracks  are mostly  generated  along  the  grain  boundaries and these initiate and propagate up to approximately 50 min after  the  sample starts being exposed to oxygen atmosphere.  Once  cracks  propagate  at  the  surface  along  the  grain  boundaries,  the  oxide  growth rate accelerates. This was evident during the HT-ESEM experiment as the  corrections required to maintain the surface of  interest at  the set working distance  and therefore with the right focus were more important. This is due to a considerable  volume  expansion. The oxidation at  this  stage becomes  severe  and the  typical  Maltese cross of  the oxide develops (see Figs. 2, 3).  When  the  oxide  growth  is  severe,  it  is  accompanied  by  shrinkage  of  the  unoxidised  carbide  core  as  revealed  in  Fig. 2  (bottom).This  picture  shows  specimens  (Set B)  oxidized  at  1273 K in  air  in  a  chamber  lift  furnace with  quenching at a certain time: specimens in the top row show their oxide layer still  intact while specimens at the bottom row have had their oxide layer removed so that  the evolution of  the unoxidised carbide core can be followed. For oxidation times  512  Oxid Met (2017) 88:509-519  123  \\x0c', 'Oxid Met (2017) 88:509-519  513  Fig. 1  a SEI of ZrC sample (Set C) exposed to 2 mbar of oxygen at 1073 K after 5 min of oxygen flow.  b SEI of ZrC exposed to 2 mbar of oxygen at 1073 K after 50 min of oxygen flow: crack formation along  the boundaries  Fig. 2  Image of ZrC specimens (Set B) oxidised at 1273 K in a chamber lift furnace: specimens at  the  top show the evolution of  the oxide layer with time,  specimens at  the bottom had their oxide layer  removed so that  the evolution of  the unoxidised carbide core can be seen  Fig. 3  SEI  of  a ZrC sample  (Set C)  oxidised  at  1073 K for  a  40 min  and b 480 min  and  high  magnification SE image of  the area highlighted with a white box in Fig. 3b c details on the monitored  corner of  the sample (dotted segments),  the initial delamination followed by spallation at  the corner  (arrow),  the appearance of  large cracks at  the corner  (a)  (white lines) and from the corner  towards the  inner core (c)  (white line) and microcracks  that  run parallel  to the carbide-oxide interface (c)  (yellow  lines) (Color figure online)  123  \\x0c', '514  Oxid Met (2017) 88:509-519  over 2 h the initial cubic inner carbide core assumes a rounded then spherical shape  whilst  the oxide develops with the characteristic Maltese cross shape (Fig. 2). The  change in shape of the unoxidised carbide core from cubic to spherical suggests that  oxidation proceeds quicker at sample corners,  therefore the monitoring of a sample  corner  in the HT-ESEM is performed (see Figs. 3, 4).  In Figs. 3 and 4, we report  the monitoring during oxidation at 2 mbar O2 and  1073 K at  the corners of  two Set C samples. These two experiments allowed the  identification of three different steps/phenomena occurring during the oxidation of  ZrC and which contribute to the formation of the Maltese cross shape for the oxide.  The first step/phenomenon is  the delamination and spallation of  the edges and  corners of the oxidized sample seen in Fig. 3a) (indicated with an arrow) but much  better evidenced in Fig. 4b). This occurs before the centres of the faces, monitored  in Fig. 1, gets heavily cracked and is thus still compact (Fig. 1a). This indicates that  the corner is the first area to fail already suggesting the oxidation process is faster at  Fig. 4  Sequence of SEI on Set C sample oxidized at 1073 K in a 2 mbar oxygen atmosphere, a sample at  the beginning of the oxidation: time 0; b delamination of the oxide layer at edges and spalling off; c crack  formation and opening up of corners; d cracks development  towards the inner core (highlighted with a  white line) and microcrack propagation  123  \\x0c', 'corners. Spallation of  the first  layer of  the oxide leaves the surfaces at  the corner  exposed to faster oxide penetration. The ensuing rapid stress build-up at the corners  is presumably responsible for  the following steps.  The second phenomenon involves crack formation at  the corner  (white lines in  Fig. 3a) and from the corner towards the inner core (crack pointed by the white lines  in Figs. 3c, 4d)  thus the opening of  the corners and edges due to the build-up of  stresses. The cracking of the corners offers a route for the oxygen to reach the inner  unoxidised ZrC core promoting the characteristic Maltese cross shape.  The development of microcracks  at  the  corners parallel  to the oxide/carbide  interface,  as  evidenced  in Fig. 3c  (yellow lines),  constitute  the  third  and  last  observed  phenomenon.  For  experiments  performed  in  air  atmosphere  they  eventually become visible by the naked eye as can be observed in Fig. 2 (samples  on the top oxidised at 120, 240 and 360 min). The formation of these microcracks  obviously further increases the accessible reactive area in the initial corner regions,  therefore accelerating drastically the oxidation reaction and the associated volume  increase at  the corners/edges compared to the face centres. Proof of  the volume  expansion occurring because of the two cracking phenomenon is indirectly given in  Fig. 3c: when the choice is made to keep the upper part  in focus, the bottom part of  the surface (bottom of Fig. 3c) gets blurred thus out of focus and this is obviously  due to a drastic change in height between the two separated parts of  the original  corner.  Another  interesting  observation was made  about  the  cracks  parallel  to  the  oxide/carbide interface: these microcracks occur in a repetitive pattern as the distance between each layer (yellow lines in Fig. 3c) is approximately 20 lm thick.  This repetitive pattern suggests that  these cracks are generated cyclically and they  are due  to debonding of  the  interface between the  carbide  and the oxide. The  carbide/oxide  interface  characterization is  also part of  the present  report  and is  discussed in the  following section. Thanks  to the use of  the HT-ESEM it was  possible  to  show that  cracking  is  not  due  to  stress  build-up  related  to  phase  transformation during cooling but instead is due to the volume expansion associated  with the oxidation of  the carbide.  Oxide Layer and Interface Characterisation  Characterisation of  the oxide layer and the carbide/oxide interface was performed  on a polished sample’s cross section annealed at 1073 K in air atmosphere for 1 h  (Set A). Figure 5 shows an SEI of the interface between the carbide and the porous  oxide layer. The intermediate region appears to be a dense pore-free layer of about 20 lm thickness. EDX analysis on this intermediate region shows the presence of  carbon, oxygen and zirconium agreeing with the suggestion by Shimada et al. of an  oxycarbide [33] compound. Nanoscale analysis on the interface was done on a FIB  milled TEM sample  cut  out  from this  region  (see Fig. 5b). Bright-field TEM  imaging (Fig. 6a)  reveals that  the compact  interface region seen by SEM (Fig. 5a,  b) is made of three zones, numbered 1, 2 and 3 in Fig. 6a. Zone 1 is indexed as the  [1 1 0] zone axis of ZrC and the SAD pattern from the circled area of Zone 1 is  shown in Fig. 6b. Zone 3 SAD pattern reveals that this region is polycrystalline with  Oxid Met (2017) 88:509-519  515  123  \\x0c', 'SAD rings of Fig. 6c matching monoclinic ZrO2. Zone 2 was not characterised via  this TEM as the smallest selected area aperture available was bigger than the area of  interest. HRTEM characterisation of the Zone 2 region seen in Fig. 6a is shown in  Fig. 7a.  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},{
  "_id": 290,
  "PDF": "ZrB2–SiC composite parts in oxyacetylenic torch tests- Experimental and computational assessment of chemical, thermal and mechanical behavior.pdf",
  "Text": "['Materials Science and Engineering A 528 (2011) 6896- 6906  Contents lists available at ScienceDirect  Materials Science and Engineering A  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m s e a  ZrB2-SiC composite parts in oxyacetylenic torch tests: Experimental and computational assessment of chemical, thermal and mechanical behavior  X.G. Wang, J. Liu ∗ , L.M. Peng ∗  CAS Key Laboratory for Mechanical Behavior and Design of Materials, Department of Modern Mechanics, School of Engineering Science, University of Science  and Technology of China, Hefei 230026, PR China  a  r  t  i  c  l  e  i  n  f  o  a  b  s  t  r  a  c  t  Article history:  Received 11 April 2011 Received in revised form 13 May 2011 Accepted 6 June 2011 Available online 13 June 2011  Keywords:  ZrB2 -SiC composite Thermal shock/ablation Geometrical factor Temperature gradient Thermal stress  1.   Introduction  The thermal shock and ablative behavior of ZrB2 -SiC pyramidal mockups with different geometrical sizes were investigated using an oxyacetylene torch. Insignificant weight or configurational changes and the absence of surface cracks after testing were observed. The excellent resistance to thermal shock and ablation is attributed to the SiC addition. A satisfactory agreement between the measured temperature distributions and FEM computations was achieved. Results indicated that both the side length and radius of nose curvature exerted a prominent effect on the temperature gradient and thermal stress magnitude inside and/or on the surface of the pyramidal mockups. The magnitudes of heat flux and boundary layer temperature, Te influenced the head temperature, T1 and the undersurface temperature, T2 in a different way. A stronger heat flux yielded higher T1 value whereas T2 remained almost constant. Comparatively, Te showed a much more remarkable effect on T2 than  it did on T1 owing to the different heat transfer mechanisms of these two parts.  © 2011 Elsevier B.V. All rights reserved.  The thermal protection materials used in hypersonic aerospace vehicles and re-usable atmosphere re-entry vehicles must be able to withstand high heat flux and high temperature. Simultaneously, they must also be endowed with good oxidation,  thermal shock resistance,  ablation  resistance  and dimensional  stability  [1-4]. Among  the ultra-high  temperature materials  (UHTMs),  carbon fiber  reinforced  carbon matrix  composite  (Cf /C) has been  considered as candidate  for  the aforementioned applications due  to its  low density, excellent high-temperature strength, high  thermal conductivity and  low coefficient of thermal expansion (CTE) [5-7]. Nevertheless, C/C composite exhibits quite poor oxidation it oxidizes above 500  C. Although an and ablation resistance as  oxidation-resistant coating may be effective to some extent in overcoming these drawbacks, cracking usually occurs during thermal cycling owing to the CTE mismatch between the C/C and coating systems [8]. As a result, the ultra-high temperature applications of C/C composite are undoubtedly restricted. Transition metal borides  including ZrB2 , HfB2 and  TiB2 ,  as another group of UHTMs, have attracted  increasing attention and interest  in recent years for their extremely high melting temper(ZrB2 3040  C, HfB2 3250  C), good  atures  thermo-chemical and  ∗ Corresponding authors. Tel.: +86 551 360 6964; fax: +86 551 360 6459. E-mail addresses: jingliu@mail.ustc.edu.cn (J. Liu), penglm@ustc.edu.cn (L.M. Peng).  0921-5093/$ - see front matter ©   2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.06.008  thermo-mechanical properties, high emissivity and thermal conductivity [9-13]. Unfortunately, monolithic ZrB2 exhibits  limited sinterability due to its covalent bonding, high melting temperature, and low self-diffusion coefficient of Zr and B. Moreover, ZrB2 is susceptible to catastrophic failure due to its inherent brittleness under thermal shock condition. Reported studies have indicated that SiC addition could enhance the densification and limit the grain growth of ZrB2 efficiently [14,15]. Other  invaluable findings demonstrate that ZrB2 -SiC composites show superior oxidation and  thermal shock resistance as compared to monolithic ZrB2 [16,17]. As major  issues  for  thermal protection applications  in severe oxidizing environments on hyposonic flight vehicles, thermal shock and ablation behavior of ZrB2 -based ultra-high temperature ceramics (UHTCs) have been investigated widely by many researchers. In general, there have been several heating approaches employed to carry out thermal shock and/or ablation tests, including oxyhydrogen torch [18], heating ovens [19,20], laser [5], plasma wind tunnel [21,22] and oxyacetylene torch [8,23]. The samples for most thermal shock experiments were heated to desired temperatures and then quenched in water or air. Under such a circumstance, the samples were subjected to a drastic temperature change for different thermal  shock  cycles usually  resulting  in  thermal  shock damage and/or catastrophic fracture. Several key parameters including residual strength, critical  temperature difference,  thermal stress and thermal shock parameters (R, R(cid:4) and R(cid:4)(cid:4) ) have been examined to describe the thermal shock behavior [19]. In previous studies, samples with simple shapes, i.e. bars for thermal shock tests [18-20], flat-face, thin column or blunt hemisphere  for ablation property                                    \\x0c', 'X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  6897  Table 1 Geometry and abbreviated designation of  rounded tips.  rectangular pyramid   specimens with  Abbreviated designation  ZS-30-3  ZS-50-3   ZS-70-3   ZS-50-1.5  Side length, L (mm)  Radius of nose curvature, r (mm) No. of cells of the body mesh   30  3  23 798   50  3  21 918   70  3  17 334   50 1.5 8 422  thermal stress distribution throughout the specimen were calculated by means of finite element method (FEM) analysis based on surface energy balance principle and a comparison was conducted with the experimentally measured results. The effects of specimen geometry, heat flux and boundary layer temperature were investigated.  2. Experimental procedures  2.1. Materials processing and sample preparation  ×  Commercially available powders were used as raw materials in  this  study. The as-received ZrB2 powders  (Dandong Chemical Co., Ltd., Dandong, China) had a mass purity of >98.5%  (with 5  <1.5 mass% B2O3 as impurity) and an average particle size of  \\u242em. SiC powders (a-phase, Shanghai Chemical Reagent Co., Shanghai, China) had a mass purity of >98.5% (impurity: <0.25% C and <0.5% 20  Fe2O3 ) and an average particle size of  \\u242em. The powder mixtures of ZrB2 -20 vol.% SiC  (hereinbelow abbreviated as ZS) was in heating oven at 80  C. Milled ball-milled  in ethanol and dried  powders were then uniaxially hot pressed in a BN-coated graphite die at 1850  C for 120 min in vacuum under 30 MPa. The bulk density and theoretical density were evaluated using the Archimedes method and the rule of mixture, respectively. The densified compacts were machined by means of electric discharge machining into three types of samples  for different tests. The bars with a size of 3 mm × 4 mm   40 mm were used for Young’s modulus evaluation using three-point bending test on a MTS 810 testing machine with a span of 30 mm and a crosshead speed of 0.5 mm/min. The flat shaped samples with a dimension of 33 mm   30 mm   3 mm were utilized  for the measurement of heat flux  from the oxyacetylene torch with different gas flow rates. In order to examine the effect of geometry on  the chemical,  thermal and mechanical behavior, several ceramic mockups with different side  length, L and radius identical vertex angle of 18  of nose curvature  radii,  r with an  were prepared by electric discharge machining (EDM). All the specimens were polished  till  the  trace of wire cutting disappeared and Table 1 presented  their abbreviated designations. For comparison, one mockup sample with L = 30 mm and r = 3.0 mm from ZrB2 -20 vol.% ZrO2 (hereinbelow abbreviated as ZZ; ZrO2 powders have a purity of 99.8% and an average particle size of 2 \\u242em provided by Yixing Ceramic Co., Yiying, China) was also prepared using the same processing route as ZZ composite.  ×  ×  2.2.  Thermal shock/ablation test and microstructural characterization  Thermal shock/ablation tests were performed in a flowing oxyacetylene torch environment. The oxyacetylene torch was obtained by  a  FeiHuan G01-30  cutting  torch gun. Gas flows of O2 and C2H2 , measured by an LZB-10 glass rotameter were 0.34 m3 /h and 0.32 m3 /h, respectively. The distance between the nozzle tip of the oxyacetylene gun and the top of the specimen was 30 mm and the inner diameter of the tip was 0.7 mm. The heat flux of the torch was measured using flat-faced ZS sample with a dimension of thickness × 30 mm  of 33 mm   3 mm. The temperature to time (T-t) curves of the front surface and back surface temperatures of the sample  ×  Fig. 1.  (a) Surface energy balance where convection, conduction and radiation were involved, (b) outside contour of specimen divided into three parts, 1, 2 and 3 and (c) mesh for calculation by FEM of ZS-50-3 specimen.  assessment [5,8,21-23] were usually used. However, the design of high-performance hyposonic vehicles generally involves relatively sharp nose tips and wing leading edges (WLEs) with the purpose of reducing the vehicle’s drag, enhancing maneuverability and performance, and improving the lift [21-24]. In this case, when the sharp nose or WLEs are heated,  in particular by means of oxyacetylene torch, thermal shock due to the temperature gradient from sharp nose to empennage part will coexist with ablation and oxidation, if the atmosphere is oxidizing. Unfortunately, there have been quite limited reports on  the shock/ablation behavior of ZrB2 -SiC with surface temperature higher than 2000  C, much less on geometrical (shape and size) effects. Therefore, the goal of this study was to perform tests on realistic parts made of ZrB2 -SiC, with  an  accurate  control of  the temperature distribution. The method was to introduce SiC particulates in ZrB2 as oxidation inhibitors and densification accelerants for hot-pressed ZrB2 -SiC composite. The densified bulk composites were machined  into sharp-nosed  rectangular pyramid with different side length and radius of nose curvature radii. Their chemical, thermal, and mechanical behaviors were investigated through exposure  to an oxyacetylene  torch flame at ultra-high  tempera(2700  C). The ablation and oxidation mechanisms were tures  revealed via microstructural examination. The  temperature and  \\x0c', '6898  X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  Fig. 2. Microstructure of the ZrB2 -20 vol.% SiC composite subjected to uniaxially hot-pressed at 1850   C for 2 h under 30 MPa.  were recorded using a Raytek pyrometer (Raynger 3i type, USA) which covers a temperature range of 600-3000  C. The test stopped after the steady-state was established by examining the T-t curves. Then, based on Fourier’s law, the heat flux was calculated to be as 1.96 MW/m2 . However, this value was,  in  fact an average one as the heat flux changes at different points on the cross-section of the oxyacetylene torch. It is documented that the heat flux of the oxyacetylene torch follows an approximate normal distribution around the radial directions. The highest heat flux in the central region took a value of about 2.08 times the calculated average value where the heat as high as 99% of the torch was concentrated [25,26]. Moreover, the radius of the specimen nose was quite small, and therefore, it is reasonable to deduce that the actual heat flux exposed to the sharp nose of the sample was 4.1 MW/m2 during the tests. During the thermal shock/ablation tests, the T-t curves on the sharp nose of  the pyramid specimens were also  recorded using the Raytek pyrometer. For the empennage part, the T-t curve was recorded by using a K-type thermocouple and an Infine TDAM 7204 16-channel data logger connected with a computer. X-ray diffraction and a PHILIPS XL-30 scanning electron microscopy (SEM) were used to characterize the chemical composition and microstructure of the sample surface after testing.  3. Modeling and computational method  The distribution of temperature and thermal stress throughout the sample during  thermal shock/ablation  tests were calculated  Fig. 4. Experimentally measured variations of temperature, T with torch exposure time, t for different samples to demonstrate the effects of side length and radius of sharp nose: (a) different L (30 mm, 50 mm and 70 mm) with the same r (3 mm) and (b) different r (3 mm and 1.5 mm) with the same L (50 mm).  by FEM on the commercial ANSYS codes. Considering the similarity in shape, the case of ZS-50-3 specimen was emphasized. Fig. 1 presents the steady-state energy balance at the surface, FEM modeling and meshing, where mutual convective and radiative heat exchanges between the sample and the boundary layer as well as  Fig. 3. Macrographs of ZS-50-1.5 specimen after exposure to oxyacetylene torch for different times.  \\x0c', 'X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  6899  Table 2 Temperature,  torch  sock/ablation tests.  exposure   time   and mass   change   obtained   from   thermal  Specimen  T1 (  C)   T2 (  C)   Time to steady state, tss (s)  Total ablation time, Ttotal (s)  \\x01m (%)  ZS-30-3 ZS-50-3  ZS-70-3  ZS-50-1.5   2630  2562  2448  2346   1209  1132  937  1199   79  88  130  64   365  345  303  290   0.10 0.005 0.06 0.002  total four  the heat conduction  inside  the sample are considered. The  unit numbers  including volumetric and surface ones  for the  meshing bodies are given in Table 1. It should be pointed out that the thermal shock/ablation process actually  involves a nonlinear transient heat transfer. The relevant material parameters, such as  thermal conductivity, specific heat and radiation, are temperature-dependent even if the heat flux load remains constant during the test. Here, the ZS ceramic is assumed to have homogeneous properties and the relevant parameters are obtained via the rule of mixture for composites. In this case, according  to energy conservation,  the  transient heat  transfer equation could be written as [C (T )]{ ˙T }  [K (T )]{T }   [Q (T,  t )],  (1)  +  =  where [C(T)] is the specific heat matrix to consider the increase of internal energy,  [K(T)]  is  the  transmitting matrix,  including conductivity, convection, emissivity and shape factor, to characterize energy expenditure,  [Q(T,t)]  is the nodal heat flux vector  including outside-absorbed and self-generated heat of  the model.  and   are the nodal temperature and its time derivative vectors, respectively. For simplicity, the above equation is re-written as [C ]{ ˙T }  [K (T )]{T }   [Q ],  {T}  { ˙T }  (2)  +  =  where  the variations of specific heat, density, and coefficient of thermal expansion with temperature are neglected and the heat flux load remains constant. The initial and boundary conditions of the relevant variables are given as follows.  (a) For T = T(r,t) (r = (x,y,z)), T0 = T(r,0) = 20  C (room temperature) as the initial condition for all cases. (b) With respect to Q:  (cid:2)  =  Q   Q1 ,  0,  ∈ /∈  r  r  (cid:3)1 (cid:3)1  ,  (cid:3)  where  is  the boundary of  the specimen,  is  the spherical area on the head and  (cid:3) 2 is the rest area of the specimen  ), respectively. Q1 (=4.1 MW/m2 ) is the initial heat  ((cid:3) 1 +   (cid:3) 2 =   (cid:3)  (cid:3) 1  flux load.  qε =  ·  (c)   −  ∈  r  (d) qi = hi ·  (T(r,   ε  t )    (T 4 (r,  is the heat flux of radiation (Te the boundary layer temperature as indicated by Fig. 1a and  the emissivity).  T 4 e ),  (cid:3)  ε  is  is  (cid:3) 2  r   t)    Te ),   (i = 1, 2, 3)  is  the heat flux of convection, neglected on the sharp nose region of the specimen compared  is  the convective heat  transfer coefficient where i (=1, 2, and 3) stands for the three regions (the individual side length ratio is 1:2:3) as shown in Fig. 1b.  to Q. hi  −  ∈  Unfortunately, as  for the convection boundary condition,  it  is difficult to determine the accurate value of hi under the present experimental conditions, where  the convection occurs between solid body and  its surrounding atmosphere. This  is attributed to the fact that hi is influenced by many complex factors such as the flowing rate of atmosphere, the temperature difference, the shape  Fig. 5. Macro-morphologies of specimens before (left in each sub-graph) and after in each sub-graph) thermal shock/ablation test: (a) ZS-30-3 (Tmax = 2630   C, (right  t = 365 s), (b) ZS-50-3 (Tmax = 2562   C, t = 345 s), (c) ZS-70-3 (Tmax = 2448   C, t = 303 s) and (d) ZS-50-1.5 (Tmax = 2346   C, t = 290 s).  and relative position of heat-transferring surface. Nevertheless, hi is assumed to take gradient values on the specimen surface. For ZS-50-3 shown  in Fig. 1b, hi takes 40, 30, 20 and 10 W/(m2 K) for regions 1, 2, 3 and the undersurface, respectively. It is worth noting that such a treatment will not produce remarkable deviations as demonstrated in previous investigations [5,27]. In fact, the convective heat transfer exerts  less  insignificant effect than conduction and radiation on  the calculated result, as  it will be addressed  in detail in the subsequent section. Then, the thermal stress, [(cid:5) ] can be calculated through  =   [M ][ε],  [(cid:5) ]   (3)  where [M] is the material elastic constant matrix including Young’s modulus and Poisson’s ratio and [ε] is the strain matrix. The interval of time,  \\x01t employed in all the calculations is automatically set by the program in the range of 0.25 s for the initial period and 3.0 s for the late stage.  4. Results and discussions  4.1. Microstructure and Young’s modulus of hot-pressed ZS composite  Fig. 2 shows a scanning electron micrograph of the polished surface of hot-pressed ZrB2 -20 vol.% SiC composite. It is evident that the dark SiC particulates are almost homogeneously distributed throughout the gray ZrB2 matrix without obvious agglomeration. The existence of SiC particles  restricted  the exaggerated grain growth of ZrB2 . However, beyond expectation, a small amount of residual pores (jet black on the SEM image) are also present, mainly due to lower sintering temperature compared to those employed in literature [19,21]. The bulk density of the sintered ZS compacts was                                                \\x0c', '6900  X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  Fig. 6.  (a) X-ray diffraction pattern of ablating products and (b) through (d) morphologies of the oxidized layer along the longitudinal surface of ZS-50-3 (marked by A, B and C regions as shown in Fig. 5b).  5.3 g/cm3 , which corresponded to a relative density of about 96% compared with the theoretical density of 5.51 g/cm3 . The Young’s modulus was estimated to be 423 GPa using the load-displacement curves obtained from three-point bending tests.  4.2. Ultra-high temperature ablation/oxidation behavior of ZS composite  Fig. 3 shows the macrographs of the ZS-50-3 sharp-nosed rectangular pyramid model at different testing times. High temperature region (bright area on the specimen) expanded rapidly from nose towards bottom, and a stable  temperature field was established after about 40 s of heating. Fig. 4 plots  the surface  temperature versus time curves of the specimens with different side  length, L and nose radius, r. Table 2 summarizes the experimentally measured results of the four specimens including steady-state surface temperature, total time and mass change after exposed to the oxyacetylene  torch with a heat flux of 4.1 MW/m2 . As evident  from Fig. 4, the surface temperature  increased drastically  in very short time, and then rose gradually, until reached a steady temperature. As a result, a large temperature gradient was generated and a high thermal stress occurred. Moreover, both the temperatures T1 of the cusp part and T2 of the empennage part decreased with  increasing L (Fig. 4a). However, T1 increased whereas T2 decreased with increasing r (cf. Fig. 4a and b). Similar to the effect of L and r on T1 , much more time, tss was needed to establish the steady-state for the ablated specimen with larger L and r. The above experimental finding can be understood  in the  following way. During  the first 10 s of heating,  large quantities of heat was absorbed by the head of specimens and cannot be transferred away or consumed in a short time due to the relatively low thermal conductivity of materials. As a result, the temperature on the head part of specimen, T1 rose rapidly. As T1 increased, obvi ous heat  loss occurred, especially  in  form of the  fourth power of temperature-dependent radiation based on Stephan-Boltzmann’s law  for blackbody radiation. Under  thus a circumstance, T1 rose gradually and eventually reached a steady value. Similarly, T2 experienced almost the same rising and stabilizing processes as T1 did. However, the empennage part gained heat mainly from the head in form of conduction and from environment in forms of conduction, convection and radiation, instead of from the high heat flux of the torch directly. Consequently, the heat transfer process needs much more time and  is extended with  increasing L. On the other hand, the heat capacity of the specimen increased with increasing L and r, implying that more heat is needed for the temperature to rise to the same level. Thus, the temperature rising rate is further lowered, and the time for steady-state establishment is extended. In our computations, specimens with identical r absorb the same power. With increasing L, the area of the cross-section and the temperature gradient increased simultaneously. In this case, according to Fourier’s law (Eq. (4)) the conductive heat flux form the cusp part to empennage part increases, and therefore, T1 decreases:  =  −k  ·  A  ˚   ·   grad(T ),  (4)  ˚  A   where   is the conductive heat flux, k is the thermal conductivity, is the cross-section area, and grad(T)  is the temperature gradient. With  increasing r, the area of specimens’ head  increased and absorbed more heat flux per unit time. As a result, higher T1 was obtained. However, as will be hereafter shown, the boundary layer temperature, Te exerted a prominent effect on T2 . Along the longitudinal direction of the oxyacetylene torch, the temperature will decrease with increasing distance away from the nozzle tip. On the other hand, the temperature at outer region of the flame is  lower than that at inner region. In this case, the average boundary layer temperature around the specimen will decrease with increasing L            \\x0c', 'X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  6901  and r as the specimen is almost totally surrounded by oxyacetylene torch, and in return T2 exhibited lower values. Fig. 5 shows the macro-morphology changes of the specimens before and after thermal shock/ablation testing.  It  is evident that no appreciable configurational changes were observed except for a small expansion on the head due to ablation/oxidation. Furthermore, no cracks were present on the surface,  indicating that the composite exhibits excellent thermal shock resistance despite the existence of large temperature gradient and thermal stress. Obvious oxidation occurred only in a small region near the head. The ablated products  identified by XRD are shown  in Fig. 6(a), indicating that the exposed surface was composed of ZrO2 , B2O3 and SiO2 . For  the ZrB2 -SiC system,  the  following oxidation processes are expected  to occur at high  temperatures  in oxidizing atmosphere [21,23,28,29]:  +  →  ZrB2 (s)    5/2O2 (g)   →   ZrO2 (s)   +   B2O3 (l)   B2O3 (l)   B2O3 (g)  SiC(s) + 3/2O2 (g)   →   SiO2 (l)   +   CO(g)   SiC(s)   SiC(s)   +  +  →  +   O2 (g)    SiO(g)    CO(g)    2SiO2 (l)   →   3SiO(g)   +   CO(g)   SiO2 (l)    SiO2 (g)   →  (5)  (6)  (7)  (8)  (9)  (10)  In general, oxidation results in a net weight gain. Nevertheless, as shown  in Table 2, only quite a small weight gain was detected for all  the specimens after ultra-high  temperature ablation. This abnormal result is mainly attributed to the exhaust of CO, SiO, SiO2 and vaporization of B2O3 and SiO2 as  indicated by reactions  (6) through (10). Fig. 6b-d displays the surface oxides’ morphology in different temperature regimes (b-ultra-high, c-intermediate, and  Fig. 7.  (a) Macro-morphologies of ZrB2 -20 wt.% ZrO2 (denoted as ZZ) specimens before and after thermal shock/ablation test (ZZ-30-3: Tmax = 2325   C, t = 326 s) and (b) micro-morphologies of the oxidized surface layer marked by A as shown in (a).  Fig. 8. Comparison between the calculated and experimentally measured temperatures  in different samples: (a) ZS-30-3, (b) ZS-50-3, (c) ZS-70-3, and (d) ZS-50-1.5. The same heat flux of 4.1 MW/m2 was imposed whereas different boundary layer temperature, Te (as shown in Fig. 6a) was chosen.  \\x0c', '6902  X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  Fig. 9. Calculated variations of temperature, T with torch exposure time, t for ZS-50-3: (a) under different heat flux while Te taking an identical value of 1100   C and (b) at different Te under an identical heat flux of 4.1 MW/m2 .  Fig. 10. Contours of temperature in ZS-50-3 specimen indicating that the steady state in temperature field was established after exposure to a heat flux of 4.1 MW/m2 for approximately 60 s.  \\x0c', 'X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  6903  Fig. 11. Contours of principal tensile thermal stress in ZS-50-3 specimen owing to the existence of temperature gradient from the sharp-nose to empennage part. It is evident that the maximum thermal stress was generated during heating about 8 s after application of a 4.1 MW/m2 heat flux.  d-low temperature) of ZS-50-3 head after ablation for 345 s under 4.1 MW/m2 . A  large number of pores with different  sizes were observed inside the ultra-high temperature region oxides. The oxidation degree became weaker with a dense oxidation layer in the regimes away from the former region. Parthasarathy et al.’s [28] and Fahrenholtz’ [29] studies demonstrated that monolithic ZrB2 experienced three regimes, depending on temperature. It is first oxidized at 800  C to yield a mixed scale of ZrO2 and B2O3 through reaction (5). Owing to the lower melting temperature (450  C),  liquid B2O3 can form a continuous  layer. In this case,  the surface  is sealed and  the  inward  transport of oxygen below 1100  C can be efficiently  inhibited. However,  liquid B2O3 begins to evaporate with increasing temperature and evolves towards outside with some pores being left inside the oxide layer. For the ZrB2 -SiC system, the oxidation resistance of ZrB2 can be  improved through the formation of a refractory silica-rich scale on the exposed surfaces which acts as an effective barrier to the inward transport of oxygen below 1500  C  [30-32]. Unfortunately, SiC cannot provide vigorous protection at ultra-high temperature, i.e. 2000  C owing to the active oxidation of SiC into SiO and the rapid evaporation of SiO2 .  In this case, ZrO2 is enriched on the surface and forms a ZrO2 skeleton where the glassy SiO2 is dispersed with the formation of a compact ZrO2 -SiO2 skeleton structure responsible  for  the oxidation resistance and configurational stability of ZrB2 -SiC specimens. It  is worth mentioning that the vaporization of B2O3 and of a small part of SiO2 on the head took away some heat and relieved thermal stress to a certain extent. Fig. 7 shows the macroand micro-morphologies of ZrB2 -20 wt.% ZrO2 mockup thermal exposure at 2325  C  sample before and after  for 326 s. Much more  significant configurational change and  severe abla Table 3 Material parameters for ZS ceramic employed in FEM computation.  Table 4 The thermal conductivity coefficient of ZS composite at different temperatures.  Cp (J/(kg K))   (cid:6)   (g/cm3 )    (10−6 /K)   ˛  576.8   5.30   6.3   (cid:7)  0.16   E (GPa)   423   ε  0.6  T (  C)   20   k (W/(m K))   48.0   200   26.7   500   21.2   1000   1500   2000   2500   3000  16.8   13.3   10.7   8.60   7.00    \\x0c', '6904  X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  Fig. 12.  (a) Typical position in the 1/8 model of the ZS-50-3 specimen and (b and c) showing the variations of temperature and thermal stress with distance from nose to empennage, respectively. L1 , L2 , L3 , and L4 are the vertical center axis, empennage center axis, mid-transverse center axis and side-surface symmetrical axis, respectively. Oi is the starting point of individual Li and si is the distance from individual Oi (i = 1, 2, 3, 4).  tion/oxidation in the head area and even in the middle part occurred at lower temperature compared to the result in Fig. 5a for ZrB2 -SiC sample with an identical geometry. In addition, the oxidized surface layer was quite loose and cracked without a skeleton microstructural feature. Accordingly, it can be inferred that ZrB2 -SiC mockup sample exhibits superior resistance to both thermal shock and ablation/oxidation.  4.3.  Calculated temperature gradient/thermal stress distribution and comparison with experimental results  The nominal physical, thermal, mechanical, and surface properties for ZS composite employed in the analyses are listed in Table 3, where the specific heat,  linear thermal expansion coefficient and Poisson’s ratio were calculated based on the rule of mixtures using the parameters of ZrB2 and SiC. The thermal conductivity coefficient of ZS composite at 300, 600, 1000, 1200 and 1500 K presented  in Table 4 were determined by Lichtenecker formula as follows [33]:  =  ln k  ϕ1 ln   k1 +   (1   −  ϕ1 )ln   k2 ,  (11)  where  ϕ1 is the volume fraction of constitute 1, k, k1 , and k2 are the thermal conductivity coefficients of composite, constitutes 1 and 2, respectively. The values of thermal conductivity coefficients at other temperatures are obtained via exponential decay fitting.  Fig. 8a-d shows the calculated variations of temperature versus torch exposure time for different specimens and comparisons were conducted with experimental data. It can be found that satisfactory agreements were achieved between the calculated and experimentally measured results (also see Table 5). It should be noted that the i.e. 1000-1200  C were selected for value of Te in a narrow band,  the four specimens. In order to confirm whether convection exerts remarkable effect, the case with convection being neglected was also analyzed. The respective temperatures at the head and undersurface of ZS-50-3 specimen were 2523 and 1106  C for the latter case, which were very close to the convection-involved values. Fig. 9a  and b  shows  the  calculated  curves  for ZS-50-3 specimen to demonstrate the influences of heat flux and Te , respecfrom 2312  C under tively. As evident  from Fig. 9a, T1 increased  to 2708  C under 5.0 MW/m2 whereas T2 3.5 MW/m2 remained almost constant as 1050  C. It  is reasonable to deduce that higher values of T1 under  larger heat flux are attributed to the  increasing heat absorbed per unit time. However, as heat loss via radiation was relatively severe, the increment of T1 was quite limited despite a large heat flux gap of 1.5 MW/m2 . The reason responsible for the unchanged value of T2 with heat flux lies in the fact that, most of the heat on undersurface was obtained from the environment instead of thermal conduction as the thermal conductivity of ZS ceramic was relatively  low. Accordingly, Te exerts significant  influence on T2 . This deduction was confirmed by the calculated variations of T-t  T-t   Table 5 Calculated (Tcal ) and experimentally measured temperatures (Texp ), length of high temperature/high thermal stress regions (LHTemp /LHStress ), and maximum thermal stress,  (cid:5) TSmax .  Specimen   Temperature (nose/empennage)   Length of high temperature and thermal stress regions   (cid:5) SSmax (MPa)  ZS-30-3 ZS-50-3  ZS-70-3 ZS-50-1.5   Tcal (  C)   2522/1147  2480/1049  2450/937  2163/1152   Texp (  C)   2630/1209  2562/1132  2448/937  2346/1199   LHTemp (mm)   LHStress (mm)   LRelative (%)  9  11  12  6   9  11  12  6   28  21  17  12   44.1 43.8 43.8 20.4        \\x0c', 'X.G. Wang et al. / Materials Science and Engineering A 528 (2011) 6896-  6906  6905  Fig. 13.   (a) Maximum thermal stress area and (b) variation of maximum thermal stress with torch exposure time.  as shown in Fig. 8b, where different values of Te were selected for ZS-50-3 specimen exposed to an identical heat flux of 4.1 MW/m2 . Difference of about 700  C in T2 values was observed when Te took values of 20 and 1000  C,  respectively. Nevertheless, T1 did not show  large difference almost without dependence upon Te as the head of the specimen was directly surrounded by the high temperature flame. Figs. 10 and 11 show  the calculated contours of  temperature and principal  tensile  thermal  stress  inside  the ZS-50-3  threedimensional body,  respectively.  It was  indicated  that while  the steady state  in temperature field was established after exposure to a heat flux of 4.1 MW/m2 for approximately 60 s, the maximum thermal stress was generated during heating about 8 s under this heat flux. Fig. 12 shows  the distribution of  temperature  (Fig. 12b) and thermal stress (Fig. 12c) with distance at typical positions of the ZS-50-3 specimen (Fig. 12a).  It can be seen that the temperature dropped drastically along the central  longitudinal, L3 and surface axis L4 , especially  in the small region near the head where  large thermal gradient existed. However, the change turned to be gentle from approximate 20 mm away from the head. Along the horizontal direction, L2 and L3 from center  to surface,  the  temperature than 10  C. As a  difference was  less  result,  large  thermal stress was generated in the corresponding high thermal gradient region and  it decreased  to a  low steady-state value at  the distance of about 20 mm away from the head. However, on the cross-section inside  the  specimen or  the undersurface, a very  small  thermal stress occurred due to the small temperature gradient. As shown in Fig. 13a, the maximum thermal stress, (cid:5) TSmax was usually generated at the tangent connecting points between the axis and the sphere head. Similarly,  (cid:5) TSmax also changed with torch exposure time (Fig. 13b). During the first few seconds of heating,  i.e. 8 s for ZS-50-3 specimen, it increased rapidly to a maximum value 48 MPa owing to the large temperature gradient. With the homogenization of temperature after 40 s of heating,  (cid:5) TSmax decreased to a steady value of 43.8 MPa. Table 5 presents the individual maximum steady state thermal stress,  four specimens.  It can be  found that  exhibited quite weak dependence on L, whereas  it was strongly affected by r as it decreased remarkably with the decrease of r. As also evident from Table 5, the lengths of both high temperature and thermal stress region decreased with decreasing L and r. However, the relative  length  increased with decreasing L and  increasing r, which implies that an adequately small radius design is beneficial to decrease  the  thermal stress, and  therefore extend  the service life of components with sharp nose tips and/or wing leading edges (WLEs).  (cid:5) TSmax of the   (cid:5) TSmax  5. Conclusions  By using tetragonal pyramid specimens with rounded tips and an oxyacetylene  torch,  the  chemical,  thermal,  and mechanical behavior of hot-pressed ZrB2 -20 vol.% SiC ultra-high-temperature ceramic mockups has been  investigated.  In general, the ZrB2 -SiC mockups exhibited almost no weight, or configurational changes after exposed  to a heat flux of 4.1 MW/m2 for up  to 365 s. No surface cracks were observed despite the  large thermal gradient and thermal stress. The ablation/oxidation resistance and configurational  stability at ultra-high  temperatures was attributed  to the  formation of a ZrO2 compact skeleton where  the oxidationyielding glassy SiO2 was dispersed. Experimental and calculated results indicated that both temperature and thermal stress distributions inside and/or on the surface of specimens were remarkably influenced by the side length and nose tip curvature radius of the pyramid. The head temperature increased with increasing heat load and was not influenced significantly by the boundary layer temperature. In contrast, the undersurface temperature showed a strong dependence on  the boundary  temperature and almost did not change with the heat flux imposed on the head. These differences were attributed to the different heat transfer mechanisms of these two parts. The aforementioned findings are beneficial to the rational design of the ultra-high temperature components with sharp nose tips and/or wing leading edges used in hypersonic aerospace vehicles and re-usable atmosphere re-entry vehicles.  Acknowledgement  This work was supported by the Program for New Century Excellent Talents in Universities (NCET-07-0790).  References  [1] K.W. Dickens, D.L. Linne, N.J. Georgiadi, AIAA 2003-0283. [2] S.P. Walder, B.J. Sullivan, AIAA 2003-6915. [3]  J. Marschall, D.C. Erlich, H. Manning, W. Duppler, D. Ellerby, M. Gasch, J. Mater. Sci. 39 (2004) 5959. [4] T.H. Squire, J. Marschall, J. Eur. Ceram. Soc. 30 (2010) 2239. [5] H.W. Li, S. Amada, M. Epinat, J. Mater. Process. Technol. 138 (2003) 423. [6] R.Y. Luo, T. Liu, J.S. Li, H.B. Zhang, Z.J. Chen, G.L. Tian, Carbon 42 (2004) 2887. [7]  J. Lachaud, Y. Aspa, G.L. Vignoles, Int. J. 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},{
  "_id": 291,
  "PDF": "ZrB2–SiC(Al) ceramics with high resistance to oxidation at 1500°C.pdf",
  "Text": "['Corrosion Science 74 (2013) 154-158  Contents lists available at SciVerse ScienceDirect  Corrosion Science  j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / c o r s c i  ZrB2-SiC(Al) ceramics with high resistance to oxidation at 1500 °C  Yiguang Wang a,⇑ , Lei Luo a,  Jing Sun a, Linan An b,⇑  a Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China b Department of Materials Science and Engineering, Advanced Materials Processing and Analysis Centre, University of Central Florida, Orlando, FL 32816, USA  a r t i c l e  i n f o  a b s t r a c t  Article history:  Received 7 December 2012 Accepted 23 April 2013 Available online 1 May 2013  Keywords:  A. Ceramic B. SEM C. Oxidation  1. Introduction  In this paper, we study the oxidation behaviour of a ZrB2/Al-doped SiC composite at 1500 °C. The composite was prepared by hot-pressing the mixture of ZrB2 and polymer-derived SiC(Al). The oxidation behaviour was studied by measuring the weight change as a function of oxidation time and by observing the structure of the oxide layer. It is shown that the ZrB2-SiC(Al) exhibits different oxidation behaviour and improved oxidation resistance as compared to the conventional ZrB2-SiC without Al-doping. The improvement in oxidation resistance is attributed to that Al-doping could increase the bond strength of the Si-O and suppress the active oxidation of SiC.  Ó 2013 Elsevier Ltd. All rights reserved.  Recently, ultrahigh temperature ceramics (UHTCs) have received extensive attentions due to their critical applications in the thermal protection systems of hypersonic aerospace vehicles [1-6]. UHTCs, including ZrB2, ZrC, HfB2, HfC, and TaC, have a set of properties that allow the use of them in harsh environments associated with the atmospheric re-entry and hypersonic flights (>3000 °C), [1,4], including high melting temperature ablation resistance, and high-mechanical strength at high temperatures. Among these UHTCs, ZrB2 is the promising one for the aerospace applications because of its lowest density (6.09 g/cm3) and good resistance to thermal shock [2,7]. Since the oxidation of ZrB2 in service environments of hypersonic flights can deteriorate the properties of the materials and could even damage the vehicles, the oxidation behaviour of ZrB2 and ZrB2-based ceramics at high temperatures is highly concerned [2,8,9]. The oxidation of ZrB2 in air leads to the formation of ZrO2 and B2O3. ZrO2 tends to form a porous structure, thus itself cannot protect the material for further oxidation. On the other hand, B2O3 is a liquid at temperatures higher than 450 °C [10]. Below 1100 °C, B2O3 forms a continuous layer to form an effective barrier to oxygen diffusion. ZrB2 thus exhibits a good oxidation resistance below 1100 °C. However, at temperatures above 1100 °C, the evaporation of B2O3 results in the lost of protection to the underlying ZrB2, leading to the paralinear oxidation kinetics between 1100-1400 °C. At even higher temperatures, the rapid evaporation of B2O3 leads to a fast linear oxidation kinetics [11]. In order to improve the  ⇑ Corresponding authors. Tel.: +86 29 88494914 (Y. Wang), tel.: +1 407 823 1009 (L. An).  E-mail addresses: wangyiguang@nwpu.edu.cn (Y. Wang), An).  lan@mail.ucf.edu (L.  0010-938X/$ see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.04.037  oxidation resistance of ZrB2 at higher temperatures, SiC was added into ZrB2 to form ZrB2-SiC composites, which can form a borosilicate layer with lower volatility than B2O3 at temperatures higher than 1200 °C [12,13]. The ZrB2-SiC thus exhibits passive oxidation behaviour over a much larger temperature range than pure ZrB2. The continuous demands for safer and faster aerospace vehicles require their leading edges and nosetips to withstand temperatures over 1500 °C [1-3]. At such high temperatures, the volatility of silica becomes obvious, leading to deterioration of the oxidation resistance of the ZrB2-SiC ceramics. Although attempts have been made to improve the oxidation resistance at temperatures higher than 1500 °C by adding other additives like TaC, MoSi2, or WC [8,14,15], little improvement was observed. Previous studies revealed that aluminium-doped polymer-derived silicon-based ceramics exhibited much better oxidation resistance than SiC ceramics without doping [16-18]. During oxidation, a unique oxide structure in which the aluminium sits in the centre of six-member Si-O rings can be formed, leading to retarding the diffusion of molecular oxygen [16,19]. Such a structure can also reinforce the strength of Si-O bonds, as indicated by a significantly improved resistance to water-vapour corrosion of the aluminium-doped SiCN ceramics [20,21]. It is known that the evaporation of silica could be considered as breaking of Si-O bonds in SiO2 to form volatile SiO. The reinforcement of Si-O bonds can lower the evaporation of silica at high temperatures. In this study, we innovatively introduce polymer-derived aluminium-doped SiC (SiC(Al)) instead of conventional SiC into ZrB2 ceramics in order to further improve the oxidation resistance of ZrB2-SiC at 1500 °C or higher. It is found that Al-doping changes the oxidation behaviour and improves the oxidation resistance of ZrB2-SiC ceramics. The oxidation mechanism of the ZrB2-SiC(Al) will be also elucidated.  \\x0c', 'Y. Wang et al. / Corrosion Science 74 (2013) 154-158  155  2. Experimental procedure  Liquid polycarbosilane (LPCS, Laboratory of Special Advanced Materials, Xiamen University, Xiamen, China) and aluminium acetylacetonate powder (purity P97%, Acros Organics, Geel, Belgium) were used as the starting materials for synthesis of SiC(Al). Previous studies [16,22] indicated that the concentration of doping aluminium should be no more than 8 atm% in order to get an optimal oxidation resistance. Accordingly, we chose the mass ration of LPCS to aluminium acetylacetonate to be 9:1. The mixture was stirred for 10 h, followed by heat-treatment at 250 °C for 2 h under the protection of flowing argon to allow the complete reaction of the two. Flourier transformation infrared (FTIR) analysis revealed that the treatment made aluminium chemically bond to the polycarbosilane to form polyaluminocarbosilane (PACS) [23]. The resultant PACS was cured at 400 °C for 4 h, followed by pyrolysis at 1000 °C for 4 h to convert to SiC(Al) ceramics. The composition of the as-synthesized SiC(Al) was measured to be SiAl0.040C1.20O0.276 [23]. For comparison, SiC powder without Al-doping was also synthesized by directly curing and pyrolyzing the LPCS using the same procedure. The obtained SiC powders had a composition of SiC1.24O0.227 [23]. For making the ceramic composites, the ZrB2 powder (0.5 lm, 99% purity, Beijing Mountain Technical Development Centre, Beijing, China) and SiC(Al) (or SiC) powder were mixed together by ball-milling for 5 h in a nylon vial with ethanol as the medium. The milling media were zirconia balls. The volume ratio of ZrB2 to SiC(Al) (or SiC) was 80 to 20. After ball-milling, the mixture was dried in a vacuum oven. The resultant mixed powders were followed by hot pressing at 1800 °C then put into a graphite die, for 1 h under a uniaxial pressure of 28 MPa in vacuum. The samples obtained by using SiC(Al) and SiC are denoted as ZrB2-SiC(Al) and ZrB2-SiC, respectively. For further comparison, the sample with aluminium concentration as the same as the ZrB2-SiC(Al) was also prepared by the same sintering process using ZrB2, polymer-derived SiC powder and AlN (1 lm, 99% purity, YiNuo High-tech Material Co., Ltd., Qinhuangdao, China) (molar ratio of 1:0.596:0.011) as the starting materials. The obtained ceramic is denoted as ZrB2-SiC-AlN. The specimens of 10 mm \\x02 5 mm \\x02 4 mm were cut from the hot-pressed materials for oxidation studies. One surface of the sample was polished to 0.5 finish using diamond paste. Isothermal oxidation study was carried out at 1500 °C in an alumina tube furnace (GSL-1600X, Hefei Kejing Materials Technology Co., Ltd., Hefei, China). Specimens were placed on zirconia plates with the polished surface upwards, and then heated to the setting temperature in the furnace with the protection of high-purity argon, followed by exposure to air flowing at 10 ml/min. After oxidation, the samples were cooled in flowing high purity argon. The weight gain as a function of oxidation time was recorded by a balance with the accuracy of 0.1 mg (Mettler Tolendo AG135, Greifensee, Switzerland). The experiments were repeated at least three times for each materials and oxidation-time intervals. The oxidized samples were characterized by X-ray diffraction (XRD, Rigaku D/max-2400, Tokyo, Japan) with Cu Ka radiation, scanning electron microscopy (SEM, JEOL 6700F, Tokyo, Japan), and energy dispersive spectroscopy (EDS).  3. Results and discussion  The densities of the obtained ceramics were measured by Archimedes method to be 5.31, 5.30 and 5.32 g/cm3 for the ZrB2-SiC, ZrB2-SiC(Al) and ZrB2-SiC-AlN, respectively, closing to their theoretical densities which were calculated by using 6.09 g/cm3 for ZrB2, 3.26 g/cm3 for AlN, and 2.6 g/cm3 for SiC(Al) [23]. The SEM  images of the polished surfaces of these ceramics are shown in Fig. 1. The dark grains in Fig. 1 are SiC and the grey ones are ZrB2. All the ceramics have the similar morphology. Fig. 2 is the XRD patterns of the ceramics. Only ZrB2 and SiC are observed in the XRD patterns; no oxides and other phases were detected from the XRD. Fig. 3 shows the plots of the specific weight change as a function of oxidation time for the three samples at 1500 °C. The plots are in the format of weight change per area vs. the square root of oxidation time, thereby the slope of the tangent of the curves is the parabolic oxidation rate constant (kp) of the materials. It is seen that the ZrB2-SiC exhibits an accelerate oxidation behaviour where the oxidation rate continually increases with time. The insert in Fig. 3 reveals that the oxidation of the ZrB2-SiC ceramic obeys linear behaviour, indicating that the oxide layer formed on ZrB2-SiC ceramic cannot provide sufficient protection to the substrate in the present case. In contrast, the ZrB2-SiC(Al) follows a typical parabolic oxidation behaviour, suggesting that the dense oxide layer was formed on the ZrB2-SiC(Al) ceramic, which can effectively protect the substrate from further oxidation. The parabolic oxidation rate constant is calculated to be 2.76 mg/cm2h1/2. The figure shows that the ZrB2-SiC-AlN also exhibits parabolic oxidation behaviour, but with the oxidation rate constant of 5.89 mg/cm2 h1/2, which is twice higher than that for the ZrB2-SiC(Al). The above results clearly demonstrate that Al-doping has profound effect on the oxidation behaviour of ZrB2-SiC based materials, including (i) Al-doping changed oxidation behaviour of the materials from linear oxidation to parabolic behaviour; and (ii) the ZrB2-SiC(Al) exhibits improved oxidation resistance than the ZrB2-SiC. Another interesting thing is that the ZrB2-SiC(Al) exhibits a lower oxidation rate than the ZrB2-SiC-AlN even though they contain the same amount of Al-doping. The high-temperature oxidation of ZrB2-SiC is a process combining weight gain due to the oxidation of ZrB2 and SiC, and weight loss due to the evaporation of B2O3 and SiO [12,24]. Therefore, direct observation of the oxide layer is intuitive to better understand its oxidation behaviour. Fig. 4 shows the cross-sections of the oxidized samples at 1500 °C for 10 h. The average thickness of the oxide layers is 390 lm, 400 lm and 60 lm for ZrB2-SiC, ZrB2- SiC-AlN and ZrB2-SiC(Al), respectively. The thickness of the oxide layer on the ZrB2-SiC(Al) is about 7 times thinner than those on the ZrB2-SiC and ZrB2-SiC-AlN, further confirming the significantly improved oxidation resistance of the ZrB2-SiC(Al). Closer examination reveals that there are three layers on the top of the ZrB2-SiC substrate (Fig. 4a and d): a very thin discontinuous porous SiO2 outer layer, a dense layer containing ZrO2-SiO2- ZrSiO4, and a SiC-depleted porous ZrB2 layer adjacent to the unaffected ZrB2-SiC substrate, similar to those reported previously [7,8,12,24]. The ZrB2-SiC-AlN has the oxide structure similar to that on the ZrB2-SiC (Fig. 4b). In contrast, the oxide layer on the top of ZrB2-SiC(Al) is drastically different (Fig. 4c). First, the oxide contains only two layers instead of three layers: a SiO2 top layer and a ZrO2-SiO2-ZrSiO4 inner layer. Second, there is no obvious SiC depleted porous ZrB2 layer. And third, the top SiO2 layer is dense and continuous. EDS analysis near to the interface between ZrB2-SiC(Al) and the inner oxide layer (Fig. 4f, Point A and Point B) confirms that there is no obvious SiC-depleted layer in the oxide. No aluminium can be detected by EDS for bother ZrB2-SiC-AlN (Fig. 4e) and ZrB2-SiC(Al). It is likely because the concentration of aluminium in the samples is lower than the limitation of EDS. The oxidation mechanism for ZrB2-SiC ceramic has been intensively studied in the past years [1,7-13,24,25]. It is well accepted that the oxidation of ZrB2-SiC at high temperatures results in three-layer structure: silica layer on the top, a middle layer of ZrSiO4 or a mixture of ZrSiO4-ZrO2-SiO2, and the third SiC-depleted layer containing ZrB2 and sometimes a little ZrO2. The third  \\x0c', '156  Y. Wang et al. / Corrosion Science 74 (2013) 154-158  Fig. 1. SEM images showing the surface morphologies of the polished ceramics (a) ZrB2-SiC; (b) ZrB2-SiC(Al); and (c) ZrB2-SiC-AlN.  Fig. 2. XRD patterns of the as-received ceramics.  layer is believed to be formed due to the active oxidation of SiC at the interface of oxide and ZrB2-SiC matrix [12,25]. According to the thermodynamic study [25,26], the oxygen partial pressure at the interface is in the range of 10\\x0011-10\\x0014 Pa at 1500 °C, low enough to allow the active oxidation of SiC [25,26]. Thus, the oxidation process of the present ZrB2-SiC can be schematically illustrated in Fig. 5. At the beginning, ZrB2 is oxidized to ZrO2 and B2O3. The B2O3 easily evaporates out at 1500 °C, leaving porous ZrO2. At the interface between the oxide and the unaffected substrate, the SiO is formed due to the active oxidation of SiC. The SiO diffuses out to the surface along porous ZrO2. With the increase in oxygen pressure at the top of the sample, SiO should be further oxidized to form SiO2 or reacted with ZrO2 to form ZrSiO4. While previous studies reported that the protective oxide layer can be formed for the ZrB2-SiC ceramic and the material exhibited parabolic oxidation behaviour at 1500 °C, discontinuous silica layer and linear oxidation behaviour are observed in the present study. It could be due to that the current oxidation experiment was performed  Fig. 3. Specific weight change as a function of square root of oxidation time for ZrB2-SiC, ZrB2-SiC(Al), and ZrB2-SiC-AlN. The insert figure shows the specific weight change as a function of oxidation time for ZrB2-SiC.  under a flowing gas condition. The transition between SiO2 and SiO can be expressed by following equation,  SiO2 ðsÞ ¼ SiOðgÞ þ 1=2O2 ðgÞ  ð1Þ  The flowing gas can bring SiO with it and make the Eq. (1) move towards right-hand direction, favouring the formation of SiO. As for the ZrB2-SiC(Al), the passive parabolic oxidation was observed. As aforementioned, the only difference between the ZrB2- SiC and ZrB2-SiC(Al) is that a small amount of aluminium existed in SiC for the ZrB2-SiC(Al) ceramics. It is thereby reasonable to deduce that aluminium doping in SiC could suppress the active oxidation of SiC. This can be understood as follow. Previous studies [20,21] demonstrated that the activity of SiO2 can be decreased by a suitable amount of Al-doping. Accordingly, Eq. (1) will move towards left-hand direction, meaning SiO2 is favoured to form over  \\x0c', 'Y. Wang et al. / Corrosion Science 74 (2013) 154-158  157  Fig. 4. Morphologies ((a) ZrB2-SiC; (b) ZrB2-SiC-AlN; (c) ZrB2-SiC(Al)) and EDS analysis ((d) ZrB2-SiC; (e) ZrB2-SiC-AlN; (f) ZrB2-SiC(Al)) of the cross-sections of the oxidized ceramics.  the dense protective SiO2 layer was formed for the  SiO. Thereby, ZrB2-SiC(Al). It is a puzzle that while the ZrB2-SiC-AlN also exhibited parabolic oxidation behaviour, its oxidation rate is higher than the ZrB2-SiC(Al) even though they contain the same amount of Al. This is likely due to that the AlN usually locates at the grain boundaries  as a sintering aid [27], instead of uniformly distributes within SiC as in polymer-derived SiC(Al). During the oxidation, the unique aluminium-doped SiO2 structure cannot be formed. 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Chou, Model of oxidation of SiC microparticles at high temperature, Corros. Sci. 50 (2008) 2367-2371. [27] F. Monteverde, A. Bellosi, Beneficial effects of AlN as sintering aid on microstructure and mechanical properties of hot-pressed ZrB2, Adv. Eng. Mater. 5 (2003) 508-512.  formation  its  J.  [22]  [23]  SiAlCN  Fig. 5. Schematic oxide structure of ZrB2-SiC ceramics.  4. Conclusions  The oxidation behaviour of the ZrB2-SiC(Al), as well as the ZrB2-SiC and ZrB2-SiC-AlN, was studied at 1500 °C for up to 10 h. The measurement on weight gain as a function of oxidation time shows that both ZrB2-SiC(Al) and ZrB2-SiC-AlN exhibit parabolic oxidation behaviour, while the ZrB2-SiC exhibits linear oxidation behaviour. The ZrB2-SiC(Al) exhibits lowest oxidation rate, Observations on the oxide layers reveal that a protective layer is formed on the ZrB2-SiC(Al), but not on the ZrB2-SiC and ZrB2- SiC-AlN. These results clearly demonstrate that the ZrB2-SiC(Al) has much improved oxidation resistance as compared to the other two. The improved oxidation resistance is ascribed to that the Si-O bonds is strengthened by Al-doping, thus the active oxidation of SiC is suppressed.  Acknowledgement  This work is financially supported by the Chinese Natural Science Foundation (Grant # 51172181 and # 90916030), ‘‘111’’ Project (B08040), and Program for New Century Excellent Talents in University.  References  [1] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000 °C + hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887-5904.  \\x0c']"
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